Harnessing Filler Materials for Enhancing Biogas Separation

Review pubs.acs.org/CR

Cite This: Chem. Rev. XXXX, XXX, XXX−XXX

Harnessing Filler Materials for Enhancing Biogas Separation Membranes Chong Yang Chuah,†,# Kunli Goh,†,‡,# Yanqin Yang,†,‡ Heqing Gong,† Wen Li,† H. Enis Karahan,†,‡ Michael D. Guiver,§,∥ Rong Wang,‡,⊥ and Tae-Hyun Bae*,†,‡ †

School of Chemical and Biomedical Engineering, Nanyang Technological University, Singapore 637459, Singapore Singapore Membrane Technology Center, Nanyang Environment and Water Research Institute, Nanyang Technological University, Singapore 637141, Singapore § State Key Laboratory of Engines, School of Mechanical Engineering, Tianjin University, Tianjin 300072, China ∥ Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), Tianjin 300072, China ⊥ School of Civil and Environmental Engineering, Nanyang Technological University, Singapore 649798, Singapore

Chem. Rev. Downloaded from pubs.acs.org by UNIV OF SOUTH DAKOTA on 08/23/18. For personal use only.

‡

S Supporting Information *

ABSTRACT: Biogas is an increasingly attractive renewable resource, envisioned to secure future energy demands and help curb global climate change. To capitalize on this resource, membrane processes and state-of-the-art membranes must efficiently recover methane (CH4) from biogas by separating carbon dioxide (CO2). Composite (a.k.a. mixed-matrix) membranes, prepared from common polymers and rationally selected/engineered fillers, are highly promising for this application. This review comprehensively examines filler materials that are capable of enhancing the CO2/CH4 separation performance of polymeric membranes. Specifically, we highlight novel synthetic strategies for engineering filler materials to develop high-performance composite membranes. Besides, as the matrix components (polymers) of composite membranes largely dictate the overall gas separation performances, we introduce a new empirical metric, the “Filler Enhancement Index” (Findex), to aid researchers in assessing the effectiveness of the fillers from a big data perspective. The Findex systematically decouples the effect of polymer matrices and critically evaluates both conventional and emerging fillers to map out a future direction for next-generation (bio)gas separation membranes. Beyond biogas separation, this review is of relevance to a broader community with interests in composite membranes for other gas separation processes, as well as water treatment applications.

CONTENTS 1. Introduction 1.1. Biogas: Composition and Upgrading 1.2. Existing Technologies for CO2/CH4 Separation 1.3. Current Status of Membrane-Based CO2/CH4 Separation 1.3.1. Overview of Membrane Processes 1.3.2. Polymeric Membranes 1.3.3. Major Challenges and Present Strategies 1.3.4. Composite (Mixed-Matrix) Membranes 2. Filler Materials 2.1. Conventional Fillers 2.1.1. Zeolites and Related Materials 2.1.2. Metal−Organic Frameworks 2.1.3. Zeolitic Imidazolate Frameworks 2.1.4. Microporous Organic Polymers 2.1.5. Carbon-Based Particles 2.1.6. Mesoporous Materials 2.2. Two-Dimensional Materials 2.2.1. Graphene-Family Materials © XXXX American Chemical Society

2.2.2. Layered (Lamellar) Silicates 2.3. One-Dimensional Materials: Carbon Nanotubes 2.3.1. Single-Walled Carbon Nanotubes 2.3.2. Multi-Walled Carbon Nanotubes 2.4. Nonporous Materials 3. Composite Membranes 3.1. Definition and Motivations 3.2. Mathematical Models of Gas Transport 3.3. Membranes and Modules 3.3.1. Membrane Fabrication Techniques 3.3.2. Module Designs 3.4. Nonideal Interfacial Morphologies 3.4.1. Origins of Defects 3.4.2. Mitigation Strategies 4. Performances of Composite Membranes 4.1. Conventional Materials 4.1.1. Zeolites and Related Materials 4.1.2. Metal−Organic Frameworks

B B C C E F I J J J J O T U W Z AA AA

AC AD AD AF AG AG AG AI AL AL AM AN AN AQ AV AV AW AX

Received: February 9, 2018

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DOI: 10.1021/acs.chemrev.8b00091 Chem. Rev. XXXX, XXX, XXX−XXX

Chemical Reviews 4.1.3. Zeolitic Imidazolate Frameworks 4.1.4. Microporous Organic Polymers 4.1.5. Carbon-Based Particles 4.1.6. Mesoporous Materials 4.2. Two-Dimensional Materials 4.2.1. Graphene-Family Materials 4.2.2. Layered (Lamellar) Silicates 4.3. One-Dimensional Materials: Carbon Nanotubes 4.4. Nonporous Materials 4.5. Process Design and Economic Analysis 4.5.1. Process Design Optimization 4.5.2. Economic Analysis 4.6. Path Toward Commercialization 5. Critical Analysis of Filler Effects 5.1. Robeson Upper Bound Limit 5.2. Filler Enhancement Index: Findex 6. Outlook and Perspectives Associated Content Supporting Information Author Information Corresponding Author ORCID Author Contributions Notes Biographies Acknowledgments List of Abbreviations References

Review

biowastes11,12 and reduces the release of CH4 (a more potent greenhouse gas than CO2) into the atmosphere (especially in the food and dairy industries).13,14 In this regard, capitalizing on biogas can potentially retard the depletion of natural gas reserves and at the same time help to fight against global climate change. To realize this potential, it necessitates upgrading of the asproduced biogas (or CO2/CH4 separation) in a robust, highperforming (in terms of separation efficiency), and costeffective way. Membrane-based separation can fulfill these requirements while offering several competitive advantages over conventional biogas upgrading processes such as amine and water scrubbing.15 However, many polymeric membranes, which dominate the gas separation industry, suffer from materials-related limitations such as physical aging, plasticization, and a commonly observed permeability-selectivity tradeoff phenomenon.16 Some earlier reviews (see refs 15, 17−19) have already addressed these limitations. Hence, our focus is geared toward the discussion of functionalized materials used as fillers in composite membranes for membrane-based biogas upgrading. To date, a wide variety of fillers including zeolites, metal−organic frameworks (MOFs), microporous organic polymers (MOPs), carbon molecular sieves (CMSs), mesoporous materials, one- and two-dimensional (1- and 2D) nanomaterials, as well as nonporous nanoparticles, have been utilized in composite membranes. Yet, few attempts have been made to critically review these fillers and their effects at a comprehensive level. In this present paper, we first provide readers with an extensive survey of both conventional and emerging fillers, emphasizing on their chemistries and innovative strategies for engineering these fillers. This is followed by a discussion on the various mathematical models, different membrane preparation methods, and types of nonidealities that occur within the composite membranes. We also review a large database comprising CO2/CH4 separation performances of composite membranes and critically evaluate it by cross-referencing results and validating against the fillers’ physicochemical properties. A new empirical metric, named “Filler Enhancement Index” (Findex), is then created to offer a reliable quantification of the filler effectiveness. Conceptually, many elements of this review can be expanded to a wider research community with interest in composite membranes for other gas separations such as natural gas, CO2 capture, as well as water treatment.

BC BC BD BF BG BG BJ BJ BM BO BO BO BQ BY BY BZ CK CM CM CM CM CM CM CM CN CN CN CP

1. INTRODUCTION Since the industrial revolution, the world’s energy consumption has escalated, driven by strong demands from the everexpanding human population and increasing pressures for economic growth.1,2 As projected by the U.S. Energy Information Administration (EIA) in The International Energy Outlook 2017 (IEO2017) report, our global energy demand will continue to expand by 28% in the next ∼25 years.3 Natural gas, which comprises primarily methane (CH4), is one of the most preferred energy resources owing to its natural abundance, robust and facile production, as well as its combustion efficiency. In addition, it has a lower carbon per unit energy and burns relatively cleaner than either coal or petroleum, making it an attractive option to evade excessive carbon dioxide (CO2) emissions.2,4,5 For this reason, the total worldwide consumption of natural gas is projected to increase by over 40% until 2040.3 However, natural gas is far from being a perfect long-term energy solution because of its nonrenewable nature and the possibility of fuel-cycle CH4 leakages, which can occur during the production, transportation, and use of natural gas.6,7 From the perspective of global warming, it is also noteworthy to emphasize that CH4 leakages, from as little as 3% of the gross U.S. natural gas production, can potentially offset the environmental benefits from replacing fossil fuels with natural gas.8 Thus, there is a need to diversify our sources of energy production. One promising approach is to harvest energy from biogas produced during anaerobic digestion (AAD) of biomass materials. Biogas is a renewable energy resource that value-adds to existing efforts in fuel-switching to natural gas.9,10 Furthermore, harvesting raw biogas provides a practical waste-to-energy solution for the elimination (and often revaluation) of

1.1. Biogas: Composition and Upgrading

Biogas is an increasingly important renewable energy resource that is produced during the AAD of organic matters in landfills, forestry, and agricultural residues, as well as industrial and municipal wastes. In 2012, global biogas production amounted to 17200 kilotonne of oil equivalent per year (ktoe y−1) and this number is forecasted to escalate to ∼33000 ktoe y−1 with a biogas market expected to reach 33.1 billion US$ by 2022.20 To cater to this demand, the number of biogas plants worldwide will grow from ∼12000 to 15000 between 2016 to 2025.21 Typically, biogas consists mainly of CH4 and CO2 with composition as summarized in Table 1. However, depending on the biomass feedstocks, the yield of biogas and its CH4 content can differ accordingly. For example, biogas produced in landfills is similar to biogas produced from organic wastes except for a higher CO2 and nitrogen (N2) content, along with a noticeably lower hydrogen sulfide (H2S) content. The B


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Table 2. General Properties of CH4, CO2, and H2S31

Table 1. Typical Compositions of As-Produced Biogas and Fossil Natural Gas16,18,22,23 composition

content

unit

biogas from organic wastes

methane, CH4 carbon dioxide, CO2 nitrogen, N2 oxygen, O2 water vapor, H2O hydrogen sulfide, H2S ammonia, NH3 siloxane, [−SiRR′O−]x calorific value

vol % vol % vol % vol % vol % ppm ppm − kJ g−1

60−70 30−40 ∼0.2 0 ∼6b 0−5000 ∼100 traces 20.2

landfill biogas

natural gasa

35−65 15−50 5−40 0−5 ∼6b 0−100 ∼5 traces 12.3

75−90 1−10 >4 0 −c 4−10000 0 0 48.0

a

property

CH4

CO2

H2S

kinetic diameter (Å) quadrupole moment (×10−26 esu cm2) dipole moment (×1018 esu cm) polarizability (×10−25 cm3) global warming potential (GWP)a atmospheric lifetime (years)

3.82 0 0 26.0 28 12

3.30 4.30 0 26.5 1 100−300

3.62 − 1.0−1.1 37.8 − 96%) when optimized by process engineering.18 For these reasons, membrane-based separation is also considered a practical and viable CO2 separation technology (see Table 3 for the merits and limitations of membrane-based separation compared to other existing CO2/CH4 separation technologies). This is evident by the increasing number of publications and patents (Figure 1) along with many reviews related to membranebased CO2 separation (see refs 25, 35−37). As such, our next section is dedicated to the discussion of the current status of membrane-based technology for CO2/CH4 separation.

presence of these incombustible gases can reduce the calorific value and potentially limit the application of biogas. In the event where high energy content is important such as for vehicle fuel or grid injection, the biogas must be upgraded (by removing CO2 and other trace contaminants) to a similar standard as natural gas (Table 1). Among the many components of biogas, the presence of trace contaminants such as H2S, NH3, siloxanes, and water vapor can reduce the market value of biogas. H2S and NH3 are corrosive in nature and can oxidize further during biogas recovery to give SO2 and NOx as precursors of acid rain.24,25 Siloxanes produce silica (SiO2) during combustion which can induce abrasion and shorten the service lifetime of engines and turbines.26 Hence, removal of these trace contaminants is necessary (see Table S1 for various methods used in H2S, NH3, and siloxane removal). To meet pipeline specifications, removal of CO2 and H2S are important. In particular, H2S is toxic and can cause a wide range of adverse health effects even at concentrations as low as 5 ppm. Above 1000 ppm, exposure to H2S is lethal. Therefore, apart from reducing the calorific value of the biogas, H2S is removed mainly due to safety concerns. From membrane-based separation perspective, the presence of H2S is undesirable as it causes corrosion of system’s components and swelling of polymer matrices.27,28 This results in a loss of performance stability of the polymeric membranes. More importantly, the cost of separation increases with increasing concentration of H2S in the feed stream, thus losing its economic advantage. For a greater discussion on H2S removal, we direct readers to a recent review by Shah and coworkers (see ref 29). On this account, the review focuses solely on the removal of CO2 from biogas for the following reasons. First, CO2 is present at a much higher concentration than H2S, which increases the corrosive nature and lowers the calorific value of the biogas to a greater extent (Table 1). Second, kinetic separation of CO2 from CH4 is easier to achieve as compared to H2S, owing to the larger difference in their kinetic diameters (Table 2). Third, CO2 is a more persistent greenhouse gas despite being less potent than CH4 (Table 2).30 Therefore, enabling a robust, highly selective, energy-efficient, and costcompetitive CO2/CH4 separation is important in securing the

1.3. Current Status of Membrane-Based CO2/CH4 Separation

The key to success lies in the customizability of membranebased CO2/CH4 separation. From fundamental physicochemical properties of the membranes to complex membrane engineering, the ability to customize greatly increases the versatility and capacity of the technology to handle different target applications, feed gases characteristics, and process operating conditions.40 For instance, novel materials such as new polymers, nanomaterials, and porous fillers facilitate the development of composite (mixed-matrix) membranes with tunable physicochemical properties to handle gas separations beyond the CO2/CH4 gas pair. In terms of process engineering, the modular nature of membrane technology also allows easy retrofitting to existing biogas plants. Besides, membrane configurations can be tailored from flat sheet to hollow fiber to deliver greater flexibility in module designs and increase membrane packing density. As a result, it becomes possible to integrate liquid absorption with membrane-based technology to realize CO2/CH4 separation using a gas−liquid contacting process. Unlike gas separation, this process, known as membrane contactor, utilizes membranes to partition a C


D

high energy input is required to desorb CO2 from the solution; it is difficult to optimize operating temperature (solubility vs reaction rate); and high regeneration temperature is needed to prevent H2S (if present) absorption

1 (ambient)

presence of H2S can induce competitive and irreversible adsorption; presence of H2O can reduce CO2 uptake; and energy intensive process as high pressure is required to drive high CO2 uptake

maximizing adsorption capacity of CO2 feasible using high surface area adsorbents and high enrichment of CH4

molecular sieves (e.g., activated carbon, zeolite, silica gels, and CMSs)

98.0

physical interaction between CO2 and adsorbent; adsorption at high pressure/desorption at low pressure 4−7

pressure swing adsorption (PSA)

high CH4 purity attainable; large quantities of biogas can be processed; heat integration can be adopted to meet efficient energy requirement; and H2S, siloxane, and other contaminants can be removed by optimizing temperature and pressure complex equipment design (e.g., compressors, heat exchangers, and turbines) is necessary for achieving harsh operating conditions (−100 °C and 40 bar)

cooling medium (for low temperature requirement)

99.0

high pressure

difference in boiling points of CO2 and CH4

cryogenic distillation

mass transfer resistance can increase if membrane is wetted; hydrophobic membrane surface prone to fouling; high energy input is required to desorb CO2 from the solution; and high regeneration temperature is needed to prevent H2S (if present) absorption

membranes: hydrophobic microporous structures (e.g., PP, PE, and PTFE)c, asymmetric membranes with dense skin layers, hydrophobic ceramic materials; liquid phases: pure water, aqueous solutions,d amines,e and amino acid salts large contact area per unit volume (small plant footprint); no foaming, channeling, flooding, and entrainment particularly at high flow rates; ease of scaling up; and flexible operation as liquid and gas phase can be independently manipulated

98.0

1 (ambient)

nonselective porous membrane as a barrier to achieve gas−liquid mass transfer without dispersing one phase into the other

membrane-based separation: gas−liquid (membrane contactor)

membrane-based separation: gas−gas

feed gas contaminants (e.g., H2O and H2S) can reduce separation efficiency; shorter lifetime of polymeric membranes as compared to other technologies; polymeric membranes possess materials-related limitations (e.g., plasticization, permeabilityselectivity trade-off); and compression may be required for feed gas with low CO2 concentration

small plant footprint; low operational and capital costs; high packing density (large membrane surface area per volume); and driving force tunable by varying transmembrane pressure

polymeric membranes; inorganic membranes; composite membranes; and facilitated-transport membranes

80.0−99.5

1 (ambient) or higher

selective barrier that separates CO2 from CH4 (based on kinetic and sorption selectivity)

a Dissolution of CO2 in water and chemical bonding to amine occur concurrently in amine scrubbing. bActivated MDEA is a mixture of MDEA and piperazine, which requires lower energy regeneration due to a lower heat of reaction as compared to using MDEA alone. cPP, polypropylene; PE, polyethylene; and PTFE, polytetrafluoroethylene. dPotential liquid adsorbents include aqueous solutions of NaOH, KOH, Na2CO3, potassium carbonate (K2CO3), sodium sulfite (Na2SO3), NaHCO3, and ammonia (NH3). eExamples of amines include MEA, DEA, MDEA, triethylamine (TEA), 2-amino-2methyl-1-propanol (AMP), diglycolamine (DGA), and diisopropanolamine (DIPA). fSolubility of CO2 in organic solvents is about five times higher than water.

high water flow rate necessary for effective water recirculation and regeneration and CO2 is soluble in water, resulting in pipeline corrosion

higher solubilityf than water scrubbing; low solvent demand reduces pumping energy; and stainless steel pipelines not essential due to noncorrosive organic solvents used regeneration of solvent is required and treatment cannot occur in the presence of H2S due to sulfur formation

limitation

high removal rate of CO2; minimal loss of CH4 (99.0

4−7

4−7

higher solubility of CO2 in water than CH4; reversible chemical bonds formation between solute and solvent

working pressure (bar) typical CH4 recovery (%) material/ solvent used

physical scrubbing; higher solubility of CO2 in solvent than CH4

physical scrubbing; higher solubility of CO2 in water than CH4

working principle

chemical absorptiona (amine scrubbing)

water scrubbing

parameter

organic scrubbing

Table 3. Comparison of the Different CO2/CH4 Separation Technologies25,32,35−37,41

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technology has a clear competitive advantage over liquid or chemical absorption processes.48 A major limitation of the membrane contactor is the higher mass transfer resistance contributed by the additional membrane resistance under a nonwetted condition (Figure 2b). If wetting occurs, this mass transfer resistance increases tremendously (Figure 2c).45 Therefore, careful deliberation on the choice of membrane materials and the nature of the membrane surfaces is essential to prevent membrane wetting. As liquid absorbents are mostly aqueous-based and hydrophilic, hydrophobic porous membranes or hydrophobic membrane surface modifications are highly desirable to minimize liquid penetration.49 Liquid absorbents can also become more corrosive with higher CO2 absorption.46 Hence, chemical stability of the membrane is another important factor to consider. By and large, PP, PE, and PTFE are known to be suitable polymeric membrane materials for membrane contactors with PFTE having higher chemical stability and greater capacity to prevent membrane wetting.45,50,51 Membrane-based gas separations, as opposed to membrane contactor processes, utilize membranes as permselective barriers to realize separation by transporting a targeted gas component faster than the other. Gas transport through membranes can occur via three working principles: solutiondiffusion, Knudsen diffusion, and facilitated transport mechanism. For solution-diffusion mechanism, the permeability of a gas is a function of its solubility and diffusivity. Solubility is a thermodynamic parameter which quantifies the amount of permeating gas molecules sorbed by the membrane under equilibrium conditions, whereas diffusivity is a kinetic parameter which measures how fast these molecules transport through the membrane. Differences in solubility arise from disparities in the molecular-level interactions between the gas molecules and membrane materials while that of diffusivity originate from differences in kinetic diameters and shapes of the gas molecules.52 Gas separation via a solution-diffusion mechanism, therefore, leverages these differences to induce dissimilar gas permeabilities and achieve separation selectivity.53 Knudsen diffusion, however, occurs when the mean free path of a gas molecule inside a membrane pore is greater than the pore size. The separation selectivity depends on the inverse square root of the molecular weight of the permeating gas components.52,53 It is noteworthy to mention that if pore sizes decrease to the order of subnanometer, such as micropores in zeolites, a sharp molecular discrimination known as “molecular sieving” is observed. Presently, Knudsen diffusion sees

Figure 1. Number of publications and patents on membrane-based CO2 separation for the last 10 years (2007−2017).38,39 *Data is calculated based on percentage of publications on membrane-based CO2 separation out of the total publications on gas separation before normalizing against data from 2008. (Figure is updated as of January 24, 2018.)

liquid phase from a gas phase, and by doing so, a high gas− liquid contact area interface at the pores of the membrane is created for efficient CO2 removal.17 These strategies, which are aimed at extending the versatility and capacity of membranebased technology for CO2/CH4 separation, will be further elaborated. 1.3.1. Overview of Membrane Processes. Membrane contactor first gained attention in the 1980’s following studies on CO2 absorption using a hollow fiber membrane contacting process.42,43 To enhance the efficiency of CO2 separation, chemical adsorption is integrated by passing a liquid adsorbent through the lumen or shell side of the hollow fibers while a feed gas (comprising CO2 and CH4) is fed to the opposite side.44 The membrane is not selective and serves only as a nondispersive barrier to separate the two phases. Instead, the selectivity arises from the liquid absorbent which has a high affinity for CO2.45 Feed gas first diffuses from the bulk gas phase to the gas−membrane interface before going through the membrane pores to reach the liquid−membrane interface. Subsequently, the absorbent liquid contacts and reacts with the CO2 in the feed gas, leading to selective CO2 mass transfer at the liquid−membrane interface into the bulk liquid phase (see working principle in Figure 2a).46,47 By leveraging membrane pores to increase the effective contact area (up to 30-fold) between the gas and liquid phases, the membrane contactor

Figure 2. Schematic illustration of (a) the working principle showing a stronger CO2 affinity by the liquid adsorbent, (b) nonwetted, and (c) wetted membranes. Nonwetted membrane (gas-filled pore) has a lower mass transfer resistance as compared to wetted membrane (liquid-filled pore). The arrows describe the cross-flow direction of the feed gas and liquid absorbent. Adapted with permission from ref 46. Copyright 2016 Elsevier. E


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Figure 3. (a) Gas permeation behavior of CO2-selective facilitated transport membrane. Adapted with permission from ref 55. Copyright 2016 Royal Society of Chemistry. (b) Different membrane morphologies such as symmetric (dense) and asymmetric (ISA and TFC) structures.

to these incentives, polymeric membranes now dominate many industrial CO2 separation applications and continue to remain highly popular in academic research.59 1.3.2.1. Membranes for Industrial Applications. The firstgeneration membranes, implemented in the 1980’s, were fabricated using cellulose acetate (CA) because of its robust performance under different CO2 feed compositions and conditions. Thereafter, a large-scale gas permeation unit that was capable of handling a CO2 flow rate of 94000 N m3 h−1 was developed by Cynara.60 In addition, several industrial processes emerged under the trade names Separax UOP,61 Envirogenics,60 and Kværner Grace Membrane System.62 It was not until 1993 that the first membrane-based biogas upgrading plant appeared in The Netherlands, with an original upgrading capacity of only 25 N m3 h−1.18 Today, this capacity has improved to 375 N m3 h−1, while bigger plants with upgrading capacities above 10000 N m3 h−1 are built to meet increasing demands.63 Different membrane configurations and polymeric materials have also been successfully commercialized for CO2 separation as shown in Table 4. For industrial

applications in the separation of binary or mixed gases such as isotope separation, hydrogen extraction, and membrane distillation.54 For facilitated transport mechanism, it emerges only for membranes with reactive carriers that can interact reversibly with one gas component but not the other. Effectively, the interacting gas diffuses through the membrane via complementary solution-diffusion and facilitated transport while the noninteracting gas transports only by solutiondiffusion (Figure 3a).55 Membrane selectivity also arises as a result of this difference. A detailed discussion of the gas transport mechanisms is given in a recent review by Galizia et al. (see ref 56). The gas transport mechanisms can be distinguished based on the structure and chemical constitution of the membranes. Typically, solution-diffusion occurs in nonporous (dense) membranes while Knudsen diffusion predominates in porous membranes, especially when the pore size is less than 0.1 μm.19,57 Nonporous membranes possess either a dense structure with an isotropic morphology or more commonly, an anisotropic structure with a thin but dense “skinlike” selective layer supported on an open porous substrate (Figure 3b). Correspondingly, the cross-sectional membrane morphology is either symmetrical or asymmetrical. Asymmetric membranes can also be further categorized into integrally skinned asymmetric (ISA) and thin film composite (TFC) membranes (Figure 3b), depending on how the selective layers are fabricated. As for membranes exhibiting facilitated transport, they usually possess reactive carriers such as amine and carboxylate functional groups, which can initiate fast reaction kinetics with the acidic CO2 molecules to transport CO2 as carbamate and bicarbonate, respectively. The CO2carrier complexes can be reversibly degenerated at the permeate side of the membrane to liberate CO2 as the product gas (Figure 3a).55,57,58 1.3.2. Polymeric Membranes. Polymeric membranes hold an advantage because they are well-established and commercially scalable. Most importantly, the fabrication process is versatile as evidenced by the extensive choice of polymeric materials (tunable permeability and separation factors), tailorable membrane morphologies and characteristics (tunable mechanical stability and permeance), membrane configurations (e.g., flat sheet and hollow fiber), and module designs. Thus, utilization of polymeric membranes for CO2 separation is compelling, especially for polymeric materials giving high CO2 permeability and selectivity membranes. Due

Table 4. Overview of Prominent Membrane Suppliers for CO2 Separations, and Their Product Portfolio at a Glance60,66 year

company

origin

polymer

1980’s

Cynara, Natco (1983) Envirogenics Grace Membrane System, Kvaerner Separex, UOP Permea, Air Products (1991) Medal, Air Liquide (1994) Ube Industries Kvæner (1998)

U.S.

CA

flat sheet/ spiral wound

U.S. U.S.

PSf

hollow fiber

France

PIa

Japan U.K.

PIa CA

MTR/ABB (2008) MTR/GKSS (2012)

U.S. Germany

Sepuran, Evonik Industries

Germany

teflon silicone rubber (PEBA) PIa

1990’s

2000’s

a

F

configuration

U.S. U.K.

flat sheet/ spiral wound hollow fiber

Information on the type of PI is not furnished. DOI: 10.1021/acs.chemrev.8b00091 Chem. Rev. XXXX, XXX, XXX−XXX

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Table 5. Comparison of CO2 Permeability and CO2/CH4 Selectivity of Different Polymeric Materialsa testing condition class PI

TR

PIM

PEO

membrane material

ratio of CO2/CH4

pressure (bar)

temp (°C)

P(CO2) (barrer)

α(CO2/CH4)

ref

6FDA-4MPD:4,4′-SDA (1:1) 6FDA-6FpDA 6FDA-DAM 6FDA-DAM:DABA (3:2) 6FDA-DAM:HAB (1:1) 6FDA-DAT 6FDA-DAT:DAM (1:1) 6FDA-durene 6FDA-durene:DABA (7:3) 6FDA-durene:DABA (9:1) 6FDA:DSDA (1:1)/4MPD:4,4′-SDA (1:1) 6FDA-IPDA 6FDA-ODA 6FDA-ODA:DAM (1:1) 6FDA-ODA:DAM (1:4) 6FDA-P1 Matrimid ODPA-P1 ODPA-TMPDA P84 PMDA-ODA Ultem allyl-TR-PBO-350 APAF-6FDA-TR450 APAF-BTDA-TR450 HPEI/450−2 HPEI/450−3 mHAB-6FDA-TR425 mHAB-6FDA-TR450 pHAB-6FDA-TR450 SPDA-SBF-TR450 6FDA-SBF-TR420 TR Ac-450 TR PrAc-450 TR Pac-450 TR HAB-4MPD-6FDA-TR TR HAB-FDA-6FDA 6FDA-DAT1 6FDA-DAT2 CoPI-TB-1 CoPI-TB-6 DPPD-6F DPPD-IMM DPPD-TMPDA PI-TB-1 PI-TB-3 PIM-1 PIM-EA-TB PIM-SBI-TB PIM-BTrip-TB PIM-BTrip-TB (166-day aging) SBFDA-DMN TPE-PIM TPE-75 TPE-25 amino-functionalized PEO/epoxy-functionalized PEO Pebax-1657b Pebax-2533b PEGDA-TRIS-A

− − 1:1 1:9 − − − − − 1:1 − − − − − 1:1 1:1 1:1 − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − − −

3.0 3.4 2.0 4.4 10.0 1.0 1.0 1.0 10.0 10.0 3.0 3.4 10.3 10.0 10.0 2.0 1.0 2.0 2.0 10.0 10.0 10.0 1.0 10.0 10.0 1.0 1.0 1.0 1.0 1.0 2.0 2.0 10.0 10.0 10.0 10.0 10.0 2.0 2.0 1.0 1.0 3.0 3.0 3.0 1.0 1.0 1.0 1.0 1.0 1.0 1.0 2.0 2.0 2.0 2.0 10.0 1.0 1.0 15.0

35 35 35 35 35 35 35 35 35 35 35 35 35 35 35 35 40 35 35 30 25 35 30 35 35 25 25 30 30 30 35 35 35 35 35 35 35 35 35 35 35 30 30 30 35 35 25 25 25 25 25 35 35 35 35 35 35 35 35

134.0 83.4 681.0 17.1 54.1 55.8 191.0 626.0 256.0 305.0 57.9 83.4 16.4 54.1 130.0 3.3 10.2 1.18 50.0 0.6 87.6 1.5 109.6 1993.0 149.0 118.0 486.0 160.0 720.0 240.0 1280.0 1160.0 320.0 530.0 500.0 460.0 249.0 120.0 210.0 158.0 330.0 261.0 392.0 1600.0 457.0 218.0 5120.0 7140.0 2900.0 13200.0 4150.0 3049.0 862.0 977.0 5203.0 180.0 120.0 351.0 715.0

30.2 41.3 21.4 34.0 18.0 50.1 31.3 18.0 19.5 13.8 35.1 41.3 49.7 23.5 23.2 25.2 33.6 41.5 30.0 5.0 12.8 39.5 49.5 17.3 38.0 30.4 28.6 30.0 23.0 31.0 15.1 20.7 26.7 18.9 19.2 25.6 19.2 37.5 29.6 23.0 17.0 20.1 16.3 14.8 17.0 32.7 15.1 10.0 6.4 9.2 14.7 15.3 20.9 16.0 12.5 6.8 19.0 8.3 7.6

69 70 71 72 73 74 74 74 75 75 69 70 76 77 77 78 79 78 80 81 82 77 83 84 84 85 85 86 86 86 87 87 88 88 88 89 89 90 90 91 91 92 92 92 93 94 95 96 96 97 97 98 99 99 99 100 101 102 103

G


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Table 5. continued testing condition class

others

membrane material PEO-PBTP PVC-g-POEM XLPEO 6FPAI aBPDA-P1 BPPOdm BPPOdp CAb DBzPBI-BuI DMPBI-BuI EVA PAET PB PBI PBI-BuI PBNPI PC PDMS PEIm PES PMP POP PPO PPEES PSfb PTMSP PVAc PVC PVDF SEBS SPEEK

ratio of CO2/CH4

pressure (bar)

temp (°C)

P(CO2) (barrer)

α(CO2/CH4)

ref

− − 1:1 − 1:1 − − 1:1 − − − − − − − − − − − − − − − − 1:1 − 1:9 − − − 3:7

0.5 1.0 1.0 4.0 2.0 1.6 1.6 1.0 20.0 20.0 4.0 2.0 4.0 20.0 20.0 2.0 3.7 1.1 3.0 1.0 2.0 − − 1.0 3.0 2.0 2.8 2.0 5.0 0.1 1.5

25 35 25 35 35 22 25 35 35 35 25 25 35 25 35 26 25 37 25 35 30 25 35 30 35 25 35 25 25 35 25

150.0 70.0 450.0 52.7 2.6 85.0 65.0 5.2 25.8 3.8 19.0 1.4 40.0 0.03 2.3 2.6 8.8 3020.0 3.8 2.6 98.7 1.8 61.0 8.0 5.5 30000.0 2.2 0.2 0.9 170.6 500.0

18.0 13.7 15.0 29.3 42.0 15.2 15.6 17.3 15.9 47.2 5.5 17.9 8.0 6.0 57.0 3.7 23.6 3.6 27.0 30.5 8.7 18.0 14.2 26.7 25.0 2.3 33.5 45.0 21.2 4.3 25.0

104 105 106 107 78 108 109 74 110 110 111 112 113 114 110 115 116 117 118 119 120 121 122 123 124 125 126 115 127 128 129

a

Membranes reported do not include blended polymer. bSeparation performance of commercial membranes are based on data in the literature and not specifications from membrane suppliers in Table 4.

applications, the choice between hollow fiber and flat sheet membrane configuration stems from several considerations including membrane surface area per unit volume of the module, feasibility of the membrane materials to be spun into hollow fibers, and most importantly, manufacturing cost.57,64 Furthermore, polymers used in commercial membranes for CO2 separation are restricted to only five to seven types of materials (Table 4) owing to practical reasons such as cost, scalability, stability, robust performance, and long-term sustainability.65 Contrastingly, many more polymeric materials with high CO2 permeability are widely studied in academic research (Table 5) given the promise of these materials in increasing the CO2 separation efficiency and harnessing the full energy potential of biogas. 1.3.2.2. Membranes for Academic Research. Current stateof-the-art membrane polymers can be classified into five main classes, namely, polyimide (PI, particularly 6FDA-based polymers), thermally rearranged (TR) polymer, polyacetylene substitutes, polymer of intrinsic microporosity (PIM), and poly(ethylene oxide)-based (PEO) polymers.55,67,68 As literature on polyacetylenes is scarce, we only compared the CO2/ CH4 membrane separation performances of the remaining four classes with reference to some other representative polymeric materials (Table 5).

Generally, membranes from state-of-the-art polymeric materials exhibit higher CO2 permeability as compared to commercially available polymers. In particular, TR polymers and PIMs have a high potential for CO2 separation membranes because they possess large free volume elements that increase diffusion coefficients of the permeating gas molecules (Table 5).130 The large free volume element of TR polymers is a result of high-temperature treatments (>350 °C), which induce thermal cyclization of the functionalized PI groups to give rigid polymer backbones such as PBI and PBO that are often insoluble. For PIMs, which are generally soluble, the contorted shapes at sites of contortion or spirocenters in the polymer structures disrupt chain packing and thus increase free volume elements.55 Resultantly, membranes derived from TR polymers and PIMs give CO2 permeability that can reach 2 orders of magnitude higher than that of commercial polymers (Table 5). Although there is a current interest in the commercial development of these membranes, it is not yet achieved owing to mechanical stability issues and physical aging effects, which result in loss of CO2 permeability and concomitant gain in permselectivity over time.56,131 Isotropic membranes with dense morphologies are usually used for academic studies on polymer properties. It is also a common feat to prepare anisotropic membranes such as ISA and TFC. The aim is to capitalize on the asymmetric H


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Table 6. Plasticization Pressure and Its CO2 Permeability of Commonly Reported Polymers testing conditiona polymer

Tg (°C)

ratio of CO2/CH4

temp (°C)

plasticization pressure, Ppl (bar)b

P(CO2) at Ppl (barrer)

ref

6FDA-6FpDA 6FDA-DAM 6FDA-DAM:DABA (3:2) 6FDA-DAT 6FDA-DAT:DAM (1:1) 6FDA-durene BPA-PC BPZ-PC CA CTA Matrimid PEIm PES PPO PSf P84 TMBPA-PC

295 393 365 319 372 422 151 175 187 185 313 199 222 210 182 300 193

1:1 1:1 1:9 1:1 1:1 1:1 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2 pure CO2

35 35 35 35 35 35 25 25 27 24 22 21 21 25 23 23 25

14 15 4 >30 15 10 31 24 11 10 12 28 27 14 34 22 13

61.0 600.0 17.0 −c 160.0 850.0 4.7 1.0 6.0 7.3 4.8 0.8 2.6 82.0 3.6 0.9 13.0

147 74 72 74 74 74 142 142 142 142 142 142 142 142 142 142 142

a The applied pressure is based on the upstream feed pressure of CO2 used. bPpl is estimated from the minimum point of a CO2 permeability against CO2 feed pressure graphical plot. cNo minimum point observed up until 30 bar.

open membrane structure, which inevitably causes a drop in the membrane selectivity.135 For this reason, membranes with high permeability have low selectivity and vice versa in a majority of instances. The Robeson plots illustrate this tradeoff by summing up empirical data into log−log plots of permeability (of the more permeable gas) against ideal selectivity.136,137 A concept called the “upper bound” emerges, whereby an overarching line defines a boundary below which all these data points exist. This boundary also outlines the state-of-the-art polymeric membranes and is observed by a multitude of other gas pairs besides CO2/CH4. Second, many gas separation membranes are derived from glassy polymers.138 Glassy polymers are more susceptible to physical aging effects as the polymer chains are in nonequilibrium states when they phase inverse below the glass transition temperature (Tg).139 This means a larger free volume than that of equilibrium states is formed between the polymer chains. Over time, as the membranes age, the polymer chains reorganize toward equilibrium and the excess free volume gets reduced, causing a decrease in CO2 permeability and an increase in CO2/CH4 selectivity.131 Third, polymer chains have certain degrees of flexibility and mobility. When the concentration of a sorbing gas inside a polymer increases, the polymer swells and initiates an expansion of the free volume to increase gas diffusivity and decrease diffusion selectivity. This phenomenon, known as plasticization, occurs at high pressures and often results in higher gas flux but lower mixed-gas selectivity.140 It is of particular concern for CO2/CH4 separation as CO2 is wellknown for its plasticizing effects and can affect a wide variety of glassy polymers from plasticization pressures as low as 4 bar (Table 6).141,142 Present strategies to address the aforementioned challenges involve the use of modified polymeric materials and postsynthetic membrane modifications to improve existing polymeric membranes. Modification of polymeric materials are usually carried out by block copolymerization, polymer blending, and chemical grafting to provide good resistance toward membrane swelling and plasticization.143−145 On the other hand, chemical cross-linking is a common modification

morphologies in conjunction with state-of-the-art polymeric materials to produce high-performance flat sheet membranes. At present, ISA membranes with selective layers of less than 0.05 μm have been realized at the laboratory scale to achieve CO2 flux that is higher than that of isotropic membranes.57 In addition, TFC (or multilayered) membranes, which consist of two or more layers of different materials, are deemed more competitive as they allow an inexpensive porous substrate to be coated with a more expensive but higher performing thin selective layer.132 The TFC design is therefore useful in reducing membrane material costs, especially for custom-made materials, such as 6FDA-based, TR polymers, and PIMs, which prices can cost more than 10 to 20 US$ g−1.57 As envisioned by Baker, custom-made polymers will gradually replace conventional polymers as selective membrane materials, leading to the eventual adoption of TFC membranes over ISA membranes in the future.65 For hollow fiber membranes, dual-layer hollow fiber membranes hold the same promise as TFC flat sheet membranes. This is exemplified by the work of Li and coworkers,133 who successfully fabricated a 6FDA-durene:1,3phenylenediamine (mPDA) (1:1)/PES dual-layer hollow fiber membranes for O2/N2 separation. The dual-layer design enabled a more expensive but better-performing 6FDAdurene:mPDA selective layer to form over a relatively cheaper PES microporous inner layer. More importantly, the fibers showed a similar O2/N2 selectivity as the intrinsic ideal selectivity of 6FDA-durene:mPDA.133 In principle, such duallayer hollow fiber membranes can be expanded to the CO2/ CH4 gas pair and, together with flat sheet TFC membranes, offer compelling opportunities for an economically competitive yet high-performance CO2/CH4 separation. 1.3.3. Major Challenges and Present Strategies. Although polymeric membranes appear promising, they suffer from several well-known materials-related limitations.132 First, most polymeric membranes suffer from a recurrent permeability-selectivity trade-off as many of them have dense membrane structures with solution-diffusion as the dominant transport mechanism (Figure 3b).134 Accordingly, permeability increases as a result of faster gas diffusivity through a more I


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technique to reduce chain mobility after primary membrane formation.72,146,147 The cross-linked membranes tend to develop higher resistance to CO2-induced plasticization and physical aging owing to greater chain rigidity.72,147 Beyond polymeric membranes, inorganic membranes have also been explored as an alternative solution considering that the materials can generally give up to 5- to 10-fold higher permselectivities than that of conventional polymeric materials.132 Furthermore, a number of them exhibit excellent chemical and thermal stabilities in addition to swelling resistance.148 A multitude of inorganic materials, including CMSs, zeolites, MOFs, and other various oxides (e.g., alumina, titania, and zirconia), can be used to prepare membranes with porous or nonporous (dense) morphologies that are suitable for CO2/CH4 separation.56,149,150 However, the cost per unit area of commercial inorganic membranes is easily a few orders of magnitude higher than that of organic polymeric membranes.132 To date, inorganic membranes have yet to be economically attested even though their scalability is hypothetically feasible.151,152 This is also exacerbated by their poor mechanical properties, which make them unfavorable for largescale implementation. For a more comprehensive discussion of the challenges and solution-focused strategies, readers are referred to earlier reviews and book chapters (see refs 56, 57, 64, and 132). Given these drawbacks, there is a strategic focus on composite or mixed-matrix membranes (MMMs) as a potential solution. The key of composite membranes lies in the capacity to engineer the transport behavior of the polymer matrix by a simple concept of dispersing a filler into a polymer phase. Fabrication of composite membranes, unlike the inorganic ones, can rely on existing polymeric materials and membrane processing techniques, which are already technically well-established and economically optimized to a large extent.132 Therefore, composite membranes offer the potential combination of high processability of polymeric materials with superior gas separation performances, thus transcending the upper bound limit without significant economic penalty. 1.3.4. Composite (Mixed-Matrix) Membranes. Composite membranes are defined when a filler material (solid phase) is integrated into a continuous polymer matrix.135 The motivation is to capitalize on the micropores of filler materials (such as molecular sieves) for size-selective gas separations. However, unlike the polymeric counterparts, filler materials lack the processability and mechanical stability to realize largearea and defect-free membranes. Exploiting them in composite structures, therefore, help to strike a balance between polymeric and inorganic membranes by taking advantage of the merits of both classes of materials (Figure 4).40 Moreover, there is a wide selection of filler materials with tunable physicochemical properties, which allows considerable versatility in membrane designs and rooms for innovative adaptations. A well-designed composite membrane takes into consideration several factors. Notably, the filler material must be uniformly distributed within the matrix, has an infinitesimal polymer−filler interfacial gap, and exhibits appropriate intrinsic porosity, chemical constitution, and physical dimension. These factors are often interconnected and may work in tandem to influence many critical aspects of a composite membrane including the filler’s states of aggregation (dispersion), pore aperture, and accessibility as well as the membrane’s chemical functionality, processability, and integrity. Hence, in light of

Figure 4. Permeability-selectivity plot highlighting the working capacities of different types of membranes.136,137 Adapted with permission from ref 153. Copyright 2012 Royal Society of Chemistry.

this complexity, we approach our discussion by first examining the fundamental chemistry behind the filler materials used in CO2/CH4 composite membranes.

2. FILLER MATERIALS The rise in filler materials for composite membranes is driven mainly by active research in CO2 capture where conventional materials such as zeolites, MOFs, CMSs, and silica nanoparticles predominate. As materials science advances, other novel materials with different functionalities and dimensionalities emerge.154 In particular, 2D graphene-family materials (GFMs) and layered silicate along with 1D carbon nanotubes (CNTs) have recently attracted widespread attention owing to their unique structural features. As such, our upcoming sections are discussed according to conventional, 2- and 1D materials as well as nonporous particles (Figure 5). 2.1. Conventional Fillers

2.1.1. Zeolites and Related Materials. Zeolites are classified as microporous crystalline aluminosilicates that are built via interconnected SiO4 and AlO4 tetrahedra as primary building units. Such units can be infinitely extended via sharing of oxygen atoms to form a three-dimensional (3D) porous structure. Due to the valence states of silicon (Si) and aluminum (Al), incorporating AlO4 tetrahedra eventually leads to an overall negative charge on the framework. As a result, the framework needs to accommodate extra framework cations such as Li+, Na+, K+, Ca2+, and Mg2+ for charge balancing purposes. Occasionally, transition metal ions such as Cu2+ and Ag+ are used to maintain charge neutrality, although 100% ion exchange is unattainable.161 The overall chemical formula of zeolite is expressed as Mx/m[(AlO2)x(SiO2)y]·zH2O, where M represents the cation with a valency of m and x and y are integers, with y/x ranging from 1 to infinity, and z is the number of water molecules in the zeolite’s voids. When y/x → ∞, the framework becomes uncharged and constitutes pure silica, such as silicalite-1 (pure silica MFI) and ITQ-12 (ITW).162,163 On the other hand, the lower limit of the Si/Al ratio is set at 1 to avoid strong electrostatic repulsions that disfavor bonding to the adjacent AlO4 tetrahedron.164,165 Zeolites can be formed naturally as part of a constituent in the cavities of volcanic and basaltic rocks. However, most natural zeolites have a low Si/Al ratio.164 Furthermore, they do not possess the necessary purity and uniformity for most industrial applications. For these reasons, synthetic zeolites are largely J


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Figure 5. Different classes of functionalized filler materials covered in this review. Zeolite LTA. Reprinted with permission from ref 155. Copyright 2007. Elsevier. Mg-MOF-74. Reprinted from ref 33. Copyright 2012 American Chemical Society. COF-8. Reprinted from ref 156. Copyright 2017 American Chemical Society. SBA-15. Reprinted with permission from ref 157. Copyright 2016 Royal Society of Chemistry. TiO2. Reprinted from ref 158. Creative Commons license CC BY 3.0. Graphene-family materials. Reprinted with permission from ref 159. Copyright 2004 Elsevier. Layered silicate. Reprinted with permission from ref 160. Copyright 2009 Nature Publishing Group.

Figure 6. Typical examples of common zeolite frameworks.155

preferred. Today, numerous methods for large-scale production of zeolites have been developed, including hydrothermal, solvothermal, ionothermal, and solvent-free syntheses, which can yield zeolites with high crystallinity and purity.166−168 These methods also provide avenues to engineer zeolitic frameworks such as LTA and FAU (Figure 6) that are unavailable in the nature. For many years, zeolites have been utilized effectively as adsorbents for CO2/CH4 separation due to various advantages such as simple production as well as high chemical and thermal stability. Their well-defined pores range from 0.5 to 1.2 nm, which allow effective discrimination between molecules of similar kinetic diameters.169,170 Generally, the separation capabilities of zeolites are considered from both equilibrium and kinetic separations. Equilibrium separation is based on the absolute difference in the uptake amount of CO2 and CH4, considering that the required time for molecular adsorption reaches infinity at a particular temperature and pressure. Kinetic separation, on the other hand, capitalizes on the difference between the diffusion rates of CO2 and CH4 in the zeolite. 2.1.1.1. Equilibrium Separation. Equilibrium separation using zeolites is affected by several factors. First and foremost is the Si/Al ratio, which is an important factor especially for

zeolites with a low Si/Al ratio. A low Si/Al ratio implies a greater number of AlO4 tetrahedra and a higher negative charge in the zeolitic framework. To compensate for this charge increase, extra cations need to be introduced through cationic displacement to increase the concentration of framework cations. The result is a substantial enhancement in the electric field strength of the framework, making it advantageous for CO2 adsorption given the larger polarizability and quadrupole moment of CO2 than CH4 (Table 2). Accordingly, zeolites with a low Si/Al ratio usually exhibit a high CO2 uptake capability.171−173 This behavior is also demonstrated through an increase in the zero-coverage isosteric heat of adsorption of CO2 with a decrease in the Si/Al ratio of zeolites.174 Although zeolites with low Si/Al ratios show strong CO2 adsorptions (even at low partial pressures), they also suffer from a more difficult desorption process, resulting in poor regeneration. In addition, the usefulness of such zeolites for CO2 adsorption is often hampered by competitive adsorption from water vapor, noting that water molecules interact closely with the electric field of zeolites given their large dipole moments. Conversely, zeolites with high Si/Al ratios may exhibit satisfactory CO 2 regeneration, but CO2 adsorption is typically compromised. K


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Figure 7. (a) The molecular trapdoor effect as illustrated by a Cs-CHA zeolite with energy levels calculated using DFT measurements and (b) the energy states of level 1 and 2 for CO2, N2 and CH4, demonstrating the substantial decrease in ΔE for CO2 as compared to N2 and CH4. Reprinted from ref 179. Copyright 2012 American Chemical Society.

more dominant for kinetic separation than for equilibrium separation. Zeolites with pore apertures of 6-membered ring and below are inappropriate as the pores are too small with respect to the kinetic diameter of the CO2 and CH4 molecules. Similarly, medium to large-pore zeolites (10- and 12membered rings) are ineffective. This is because the diffusion rates of both CO2 and CH4 are indistinguishable by the large pores as substantiated by Cavenati et al.184 and Zhang et al.185 using zeolite 13X (a typical 12-membered ring zeolite). Therefore, resolving the pore apertures of zeolites is crucial for CO2/CH4 separation. On this account, zeolites with 8membered rings, such as LTA, CHA, ERI, and DDR, are deemed ideal for biogas separation given that the kinetic diameters of the two gases fall perfectly within the typical pore aperture ranges of such zeolites (Table 2 and Table 7).186 A

The second factor is the cation type and its position in the framework. Tuning the charge of zeolites can be realized by integrating different cations via a simple ion-exchange process. This ion-exchange process is a typical approach for optimizing the electric field strength and pore volume of zeolites for CO2 molecules.175 Furthermore, the Lewis acid−base interaction with the oxygen atom of the CO2 molecule can be strengthened using appropriate cations. Small alkali and alkali-earth metals such as Na, Ca, and Mg are common choices of monovalent and multivalent cations for ionexchange.161 In few instances, large electropositive monovalent or multivalent cations such as K and Cs are investigated as well.176−178 However, contrary to the conventional molecular sieving mechanism or physisorption, ion-exchange with a large cation can induce a “molecular trapdoor effect” where the cation migrates from its original position to create an unobstructed exclusive transport passage for CO2 molecules (Figure 7a).179−181 This effect was validated through an ab initio density functional theory (DFT) calculation where the energy barrier, ΔE, was found be lower for CO2 as compared to those for N2 and CH4 (ΔECO2 < ΔEN2 ≈ ΔECH4) (Figure 7b). Nevertheless, experimental corroboration of such “trapdoor effect” is challenging and not widely studied to date.182 The third factor involves the structure of the zeolites. Equilibrium CO2 uptakes are known to be influenced by the overall pore volume and surface area of the zeolites. One way to tune these parameters is to tailor the number of membered rings in the zeolitic framework. A study by Yang et al.183 substantiated this concept by demonstrating an increase in the surface area and pore volume of the 8-, 10-, and 12-membered ring zeolites through N2 sorption isotherms at 77 K. The CO2 uptake was reported to improve in increasing order of the membered ring, although the CO2/CH4 selectivity showed an increase in a peculiar order of the 10-, 8-, and 12-membered rings. The anomalous result for the 12-membered ring zeolite was attributed to a reduction of zeolite to nanocrystals, as well as the lower Si/Al ratio, which indicated the additional contribution of electronic interaction with the CO2 molecules.183 2.1.1.2. Kinetic Separation. While equilibrium separation is largely dependent on the affinity between adsorbent and adsorbate, kinetic separation exploits the difference in the diffusion rate of CO2 and CH4 through the pore aperture of an adsorbent. For this reason, the effect of zeolite structures is

Table 7. Effect of Different x-Membered Ring on the Maximum and Typical Pore Aperture188 x-membered ring

maximum pore aperture (nm)

typical pore aperture (nm)

4 5 6 8 10 12

0.16 0.15 0.28 0.43 0.63 0.80

− − − 0.30−0.45 0.45−0.60 0.60−0.80

study by Hudson et al.187 asserted this claim using zeolite SSZ13 (CHA) by showing strong discrimination between CO2 and CH4 molecules. Results also demonstrated a significant decrease in energy barrier as CO2 molecule passed through an 8-membered ring window but not for CH4 molecule (Figure 8). An overview of the CO2 and CH4 diffusion rates of common zeolites are provided in Table 8 and Table 9. At present, there is a limited number of studies on the diffusion rates of CO2 and CH4 in 8-membered ring zeolites. As such, only a few observations can be ascertained. For example, a molecular dynamics (MD) simulation by Sholl et al.189 revealed that a 3D pore could create a stronger steric hindrance to light gas molecules than a 1D pore. In general, the diffusivity of CO2 is faster than CH4, owing to preferential CO2 adsorption on the 8-membered ring as discussed earlier. Future research efforts should, therefore, focus on optimizing experimental designs to L


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Figure 8. Plot of potential energy (a.u.) as a function of the out of plane angle (degree) transport of (a) CO2 and (b) CH4 through an isolated 8membered ring window of SSZ-13. Reprinted from ref 187. Copyright 2012 American Chemical Society.

Table 8. Comparison between Diffusivity (D) and Activation Energy (E) of CO2 and CH4 Based on Different Zeolite Frameworks186 zeolite framework

D(CH4) (m2 s−1)

temp (K)

H-ZSM-5 H-ZSM-5 DD3R 4A (LTA) 5A (LTA)

334 250 298 300 (CH4); 500 (CO2) 300

E(CH4) (kJ mol−1)

D(CO2) (m2 s−1)

E(CO2) (kJ mol−1)

4 4 − 24 4

− − 1.0 × 10−10 3.0 × 10−12 1.5 × 10−8

− − − 23 9.4

−8

≈10 3.0 × 10−9 1.7 × 10−12 5.0 × 10−15 10−9

Table 9. Summary of Selected Zeolites That Could Be Utilized in CO2 Separation zeolite structure BEA CHA

DDR FAU

KFI LTA MFI MWW RHO

cation +

K Na+ H+ Cu2+ NH4+ − Li+ Li+ Na+ Na+ Li+ Na+ Ca2+ Na+ Na+ K+ Li+ K+

zeolite framework

crystal density (g cm−3)

pore size (nm)

CO2 uptake (mmol g−1)a

total pore volume (cm3 g−1)

ref

K-BEA Na-BEA H-SSZ-13 Cu-SSZ-13 SAPO-34 DD3R Li-X Li-Y Na-X (13X) Na-Y Li-ZK-5 Na-ZK-5 Ca-A (5A) Na-A (4A) Na-ZSM-5 K-MCM-22 Li-RHO K-RHO

− − 2.0−2.2 2.0−2.2 2.0−2.2 1.71 1.54 1.42 1.54 1.42 − − 1.48 1.52 1.76 − − −

0.60 0.60 0.38 0.38 0.38 0.36 × 0.44 0.74 0.74 0.74 0.74 0.39 0.39 0.48 0.38 0.51 × 0.55 0.40 × 0.59 0.41 0.41

3.0 2.8 3.98b 3.75 4.55 1.32 5.62 5.21 5.67 5.18 4.9b 4.0b 5.02 3.07 1.56 2.7 4.96c 4.50c

0.41 0.52 0.27 0.25 0.28 0.15 − 0.34 0.33 0.34 0.22 0.22 0.29 0.21 0.22 − 0.36 0.36

190 190 187 187 191 192 193 193 161 193 194 194 161 161 195 196 176 176

a

CO2 feed condition: 298 K and 1 bar. bCO2 feed condition: 303 K and 1 bar. cCO2 feed condition: 303 K and 0.8 bar.

framework can be readily modified by various cations, including zirconium (Zr), vanadium(V), tin (Sn), and niobium (Nb), without losing the overall coordination. Potentially, this suggests the easy exchange of various charge-equalizing cations. Furthermore, the formation of titanosilicate frameworks is not restricted to just conventional 3D porous structures. Successful formation of nanoporous layered materials such as Jilin-DavyFaraday (JDF-L1) and AMH-3 have been reported (see section 2.2.2). Today, the definition of zeotype are expanded to include zeolitic imidazolate frameworks (ZIFs) which portray a similar framework topology.199 This will be explained further in section 2.1.3. The first derivative of ZSM-5 (MFI zeolite), termed as TS-1, was produced via substitution of Si4+ to Ti4+ in an arrangement of tetrahedral TiO4 and SiO4 units (Figure 9a). Subsequent

circumvent complications arising from the presence of ionexchangeable cations and the effect of the pore dimensions on the intercrystal diffusivity. There is also a need to simplify elaborated experimental measurements and mathematical modeling. These efforts will be instrumental in reconciling the poor reproducibility of the CO2 diffusivity parameter. 2.1.1.3. Zeotypes (Zeolite-Like Materials). Zeotypes are classified as any crystalline materials that share similar topologies as zeolites. The materials were first synthesized in 1983 by doping titanium (Ti) on pure silicon dioxide (SiO2) instead of Al.197 Subsequently, crystalline Ti molecular sieves were successfully constructed based on effective adjustments of the TiO6 and SiO4 tetrahedra. These materials are termed Engelhardt Titanosilicate (ETS), with two zeolites (ETS-4 and ETS-10) being reported by Kuznicki.198 The resulting Ti M


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Figure 9. (a) TS-1 with isomorphous substitution of Si by Ti displayed as orange and green, respectively. Reprinted with permission from ref 206. Copyright 2016 Royal Society of Chemistry. (b) ETS-4 with Si, Ti, and O shown as yellow, green, and red, respectively. Reprinted with permission from ref 201. Copyright 2001 Nature Publishing Group. (c) Comparison of CO2 adsorption isotherms of divalent Ca2+, Sr2+, and Ba2+ cationexchanged ETS-4 at 25 °C with pretreatment temperatures at 100 and 200 °C. Reprinted with permission from ref 203. Copyright 2011 Elsevier.

Table 10. Overview of the Key Parameters of Representative Zeotype Titanosilicates That Show Promise in CO2 Separation titanosilicate

cation


pore size (nm)

TS-1 ETS-2 N1-ETS-2e N2-ETS-2e N3-ETS-2e ETS-4

− Na+ Na+ Na+ Na+ Ba2+ Ca2+ Ca2+/H+ Na+ Sr2+ Zn2+ Ba2+ Ba2+/H+ Na+ Ba2+ Ca2+/H+ Zn2+

− − − − − − − − 2.20 − − − − 1.93 − − −

− −d −d −d −d − − − 0.3−0.4 − − − − 0.49 × 0.76 − − −

ETS-10

RPZ

b

CO2 uptake (mmol g−1)a c

23.71 − 1.08f 0.98f 0.90f 2.20 0.67 2.30 2.89 1.34 2.10 2.30 2.00 3.12 1.20 1.90 1.80


ref

0.1 − − −d −d − − − − − − − − − − − −

200 207 207 207 207 203 203 208 209 203 208 208 208 209 208 208 208

a CO2 feed condition: 298 K and 1 bar. bNo additional metal source other than Si and Ti. cCO2 uptake remains unknown. The H2O uptake at 323 K and 1 bar was instead represented. dNonporous. eN1, (3-aminopropyl)trimethoxysilane; N2, [3-(2-aminoethylamino)propyl] trimethoxysilane; N3, 1-(3-trimethoxysilylpropyl)diethylenetriamine). fCO2 feed condition: 303 K and 1 bar.

increased, the number of sites available for adsorption increased. This was evident from a report by Mirajkar et al.,200 where the H2O equilibrium sorption of TS-1 was found to be proportionate to the Ti content in the framework. However, a continuous increase in the Ti content can be detrimental as amorphous titanium dioxide (TiO2), which reduces the pore volumes in the framework, is generated. Thus,

studies investigated the uptake capacity of TS-1 by increasing the amount of Ti4+. It was found that when the Si/Ti ratio decreased, an increase in the electric field strength (similar to the typical aluminosilicate zeolites) occurred despite both Si4+ and Ti4+ being isovalent. This was attributed to bond polarization induced by the huge difference in electronegativity between Ti and Si (1.32 vs 1.74). Hence, as the Ti content N


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dependent and/or diffusion-controlled molecular sieving mechanism.219−221 Tuning the pore size and its distribution, creating flexible frameworks, and easy access to open metal sites, as well as tailoring the physicochemical properties of open metal sites of MOFs are all important design parameters that require careful consideration. 2.1.2.1. Molecular Sieving. The reticular approach222 to synthesizing MOFs allows preferential coordination of metal ions or clusters with rigid organic linkers. Accordingly, an isoreticular series of MOF compounds with same topologies but different pore sizes and functionalities can be prepared by customizing the organic linkers. One example is the isoreticular MOFs (IRMOFs) series of [Zn4O(bdc)3] (MOF-5), which can be loaded with pendant functional groups such as −Br, −NH2, −OC3H7, −OC5H11, −C2H4, and −C4H4 from the organic linkers (Figure 10a).223 As the pendant groups get

the Si/Ti ratio needs careful optimization for efficient CO2/ CH4 separation. Other common structures with zeolite-like properties are ETS-4 and ETS-10. ETS-4 comprises a mixed octahedraltetrahedral framework (similar structure to the mineral zorite)201 and small pores (8-membered ring), ranging from 0.3 to 0.4 nm (Figure 9b).202 Compared to ETS-10, ETS-4 displays a much lower thermal stability, especially when Na+ cation is used. It has been reported that the structure comes apart at 200 °C to yield an amorphous material with a collapsed channel system.201 To rectify this problem, ionexchange with a larger divalent cation such as Mg2+, Ca2+, Sr2+, and Ba2+ is used to enhance the thermal stability of ETS-4. Park et al.203 demonstrated stable CO2 uptakes in Ca2+-, Sr2+-, and Ba2+-exchanged ETS-4 under pretreatment temperatures of 100 and 200 °C (Figure 9c). The uptake kinetics of gases can also be improved by pore widening through replacing two Na+ with one divalent cation.203 ETS-10, on the other hand, is another microporous titanosilicate, which is constructed based on the corner-sharing of TiO6 and SiO4 tetrahedra through a bridging oxygen atom.204 Characterization results show that such chains are parallel to the orthogonal channels with rods containing Ti chains being stacked in an alternate direction to form a large 12-membered ring framework. ETS-10 adsorbs CO2 favorably at low temperatures given that its basicity allows effective adsorption of the acidic gas.205 The CO2 adsorption capacities of representative zeotype titanosilicates are summarized in Table 10. 2.1.2. Metal−Organic Frameworks. Metal−organic frameworks (MOFs), which are often described as porous coordination polymers (PCPs) or porous coordination networks (PCNs), have attracted widespread attention in many fields (such as molecular separation, heterogeneous catalysis, and gas storage), owing to their unique structural properties, unprecedentedly high accessible internal surface areas (up to 7000 m2 g−1), high void volumes (up to 90%), low densities (down to 0.19 g cm−3), as well as narrow pore size distributions.210,211 By and large, MOFs are crystalline materials constructed based on coordination bonds between inorganic metal-based nodes (single ions or clusters) and organic bridging linkers (multidentate ligands). Numerous methodologies have been developed to yield highly crystalline MOFs. In particular, conventional approaches such as hydrothermal212 and solvothermal213 methods are widely adopted in the literature. The choice of solvent is often critical to the synthesis, although many solvents used as structuredirecting templates have yet to be thoroughly investigated. As the synthesis of MOFs becomes mature, other methods including sonochemical,214 microwave-assisted,215 mechanochemical (ball milling),216 as well as electrodeposition217 synthesis emerged. The alternatives not only strengthen the synthetic capability and versatility but also offer size and morphology control of the MOFs for specific applications. The various synthesis routes developed to date are summarized in an extensive review by Stock and Biswas.218 In this section, our focus is to discuss the use of MOFs to achieve effective CO2/CH4 separation by exploiting differences in the kinetic diameters and electronic properties (quadrupole moment and polarizability) of the CO2 and CH4 molecules. Generally, selective capture of CO2 stems from (1) adsorptive selectivity (equilibrium separation), which capitalizes on the strong CO2 affinity toward the pore surfaces, and (2) sizebased selectivity (kinetic separation), which is built on the size-

Figure 10. (a) Isoreticular MOFs with pore sizes ranging from 3.8 to 28.8 Å using various organic linkers. Reprinted with permission from ref 223. Copyright 2002 American Association for the Advancement of Science. (b) Structures of SIFSIX-2-Cu, SIFSIX-2-Cu-i, and SIFSIX-3-Zn. The effective pore size of SIFSIX can be tuned using short chain ligand to give SIFSIX-3-Zn with pore size of 3.84 Å. Reprinted with permission from ref 226. Copyright 2013 Nature Publishing Group.

larger and longer, the pore size of the IRMOFs experience a dramatic increase from 3.8 to 28.8 Å, giving evidence of pore size control using reticular chemistry.221 The fine-tuning of pore size can also be realized by the formation of an interpenetration network. For instance, the pore size of a noninterpenetrated MOF-5 was reduced from 8.6 to 5.3−5.7 Å after successful formation of an interpenetrated MOF-5 framework.224,225 Nevertheless, it is challenging to further decrease the pore size of MOF-5 to within a small pore window that is between the kinetic diameters of CO2 and CH4 molecule (Table 2). As such, the promise of MOF-5 for CO2/ CH4 separation by molecular sieving remains debatable. On this account, the molecular sieving effect is often coupled with electrostatic interactions to favor CO2 uptake. For example, SIFSIX-2-Cu is prepared from 4,4′-dipyridylacetylene (dpa) and SiF62− anions (SIFSIX) to afford primitive-cubic nets with square channels of pore dimension 13.05 Å (Figure 10b).226 The interpenetrated polymorph, SIFSIX-2-Cu-i, consists of doubly interpenetrated nets which are isostructural to the nets of SIFSIX-2-Cu but overlapped in a staggered manner to yield pores of dimension 5.15 Å (Figure 10b). O


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Figure 11. (a) CO2 uptake profile of MIL-53 (Cr) at 31 °C with a clear hysteresis loop observed during adsorption and desorption of CO2. (b) Powder X-ray diffraction (PXRD) of MIL-53(Cr) as a function of CO2 pressure at 20 °C with a clear shift in 2θ toward a lower angle as the CO2 applied pressure increases. Reprinted with permission from ref 230. Copyright 2007 Wiley-VCH.

One of the earliest examples of a flexible framework is the MIL-53 series. MIL-53 is constructed based on a 1,4benzenedicarboxylate (bdc) ligand with metals such as Al and chromium (Cr). The CO2 adsorption proceeds in a twostage stepwise mechanism: (1) a sharp increase at low CO2 pressure followed by a plateau up to 4 bar before (2) another sharp CO2 adsorption (Figure 11a).230,231 This observation can be explained by the expression “breathing” effect, in which the CO2 molecules (from the initial uptake) exert a strong interaction with the −OH groups on the chains of the framework to cause a shrinkage of the structure at low pressure, leading to the plateau between 1 to 4 bar. As pressure increases, more CO2 molecules are admitted into the cell to reopen the structure, prompting further adsorption.221 This “breathing” phenomenon was exemplified by a shift in the 2θ peaks toward the lower angle with an increase in the CO2 partial pressure (Figure 11b). For CH4 adsorption, no stepwise mechanism is observed due to a lack of interaction with the framework. As such, a CO2/CH4 binary gas adsorption of MIL-53(Cr) reveals a “breathing” effect similar to pure CO2 adsorption, verifying that the mechanism is controlled mainly by the partial pressure of CO2 and not CH4.232 Another study on MIL-53(Al) shows that the “breathing” effect of the framework results in a sharp decrease in the CO2/CH4 selectivity from ∼7 to ∼4 as the pressure exceeds 5 bar.233 This is due to an enlargement of the cell volume at higher pressure, resulting in unobstructed access to both CO2 and CH4 molecules. 2.1.2.3. Coordinatively Unsaturated Metal Sites. Open metal sites in some MOFs are useful in improving the CO2 affinity and CO2/CH4 selectivity. By removing the terminal solvent molecules, the metal sites are made coordinatively unsaturated, which can then act as strong Lewis acids to polarize CO2 molecules to form M···OCO adducts.234 In addition, they can be chemically grafted with relevant functional groups to enhance the CO2 capture capacity of MOFs. However, the effective use of such MOFs is only limited to applications that do not require exposure to humidity. This is because the lone pair of electrons on the oxygen atom of H2O can easily coordinate to the open metal sites, rendering them useless to CO2 capture. One example is Cu3(btc)2 (btc3− = 1,3,5-benzenetricarboxylate), which is commonly referred to as HKUST-1 or CuBTC. It has been reported to possess coordinatively unsaturated metal sites and is constructed based on multiple networks of

Despite a smaller pore size, SIFSIX-2-Cu-i exhibited higher CO2 capture capacity (5.41 mmol g−1) and CO2/CH4 selectivity (33 at 25 °C and 1 bar) as predicted by Ideal Adsorbed Solution Theory (IAST) calculations. Comparatively, SIFSIX-2-Cu displayed a CO2 uptake and selectivity of only 1.84 mmol g−1 and 5.3, respectively. The improved CO2 separation performance was attributed to the narrower pores of SIFSIX-2-Cu-i, which gave rise to a more favorable CO2 interaction driven by a better overlap of the attractive potential fields in the opposite walls. Subsequently, SIFSIX-3-Zn (an IRMOF of SIFSIX-2-Cu) was synthesized using a shorter pyrazine organic ligand to realize a pore dimension of 3.84 Å.226 Breakthrough studies revealed that SIFSIX-3-Zn had a higher CO2/CH4 selectivity of 109 as compared to 51 of SIFSIX-2-Cu-i at a CO2/CH4 feed ratio of 1:1. At a lower CO2 pressure of 10 mbar (or 1% CO2 loading), SIFSIX-3-Zn even demonstrated an order of magnitude higher volumetric CO2 uptake as compared to SIFSIX-2-Cu-i.226,227 The Cu analogue of SIFSIX-3-Zn (known as SIFSIX-3-Cu) was then constructed by replacing Zn2+ by Cu2+ to induce a geometrical distortion via a Jahn−Teller effect.227 As a result, the pore size of SIFSIX3-Cu was decreased further to produce one of the highest uptakes at a CO2 pressure as low as 0.4 mbar. These results also show that narrowing the pore size helps to reduce the distance between the CO2 molecules and F atoms of SIFSIX, thereby strengthening the interactions of adsorbed CO2 molecules in the channels. 2.1.2.2. Flexible Frameworks. As a coordination network, the MOFs discussed so far exhibit frameworks which are rigid even though some of them can be considered less stiff as compared to zeolites.228 Yet, there is a class of MOFs with flexible frameworks that undergo reversible structural changes when exposed to external stimuli such as pressure, temperature, and interaction with guest molecules. The structural flexibility stems from diverse polarities in the backbones or pendant side groups of the frameworks.229 Reversible changes are typically characterized by a distinct step (gate opening) in the CO2 adsorption isotherm at a specific pressure and a hysteresis between the adsorption and desorption branch of an isotherm (Figure 11a). When the framework is triggered to open, an enlargement in the aperture can occur by more than 5 Å, which may lead to a loss in the CO2/CH4 mixed-gas selectivity.219 Hence, flexible MOFs have limited application for CO2/CH4 separation to date. P


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Figure 12. (a) Structure of HKUST-1. Cu, C, and O atoms are represented by green, gray, and red spheres, with H atoms being omitted for clarity. Reprinted from ref 33. Copyright 2012 American Chemical Society. (b) CO2 and CH4 uptake of HKUST-1 and zeolite 13X at 25 °C. Adapted from ref 237. Copyright 2009 American Chemical Society. (c) Structure of M2(dobdc). M, C, and O atoms are represented by black, gray, and red spheres. Reprinted from ref 33. Copyright 2012 American Chemical Society. (d) CO2 uptake of isostructural M2(dobdc) at 23 °C. Adapted from ref 238. Copyright 2008 American Chemical Society.

paddlewheel Cu2(COO)4 nodes and triangular btc3− linkers (Figure 12a).235,236 Effective removal of solvent molecules bound to the axial sites of the Cu2+ metal center allows the creation of adsorption sites to give preferential adsorption of CO2 (12.7 mmol g−1) over CH4 (4.6 mmol g−1) at 25 °C and 15 bar (Figure 12b).237 Besides, it exhibits a lower isosteric heat of adsorption (30 kJ mol−1) as compared to other conventional adsorbents such as zeolite 13X (49 kJ mol−1). Therefore, only a small energy penalty is incurred when regenerating such adsorbents for subsequent CO2 adsorption. Another series of MOFs that contains open metal sites is the M2(dobdc) (M-MOF-74, M = Mg, Mn, Fe, Co, Ni, Cu, or Zn; H4(dobdc) = 2,5-dihydroxyterephthalic acid). This series of MOFs has a framework based on a honeycomb-type topology with a 1D channel of 11−12 Å in diameter that is aligned with the exposed M2+ sites (Figure 12c). Numerous isostructural frameworks of the M2(dobdc) series with different open metal sites give rise to different CO2 adsorption performances. For example, the adsorption performances of Ni2(dobdc) and Co2(dobdc) registered about 50% of Mg2(dobdc) with a CO2 uptake of 23.6 wt % (as compared to 11.6 and 11.7 wt %) at 296 K and 0.1 bar (Figure 12d).238 This is probably due to a lower CO 2 adsorption enthalpy of Ni 2 (dobdc) and Co2(dobdc) in contrast to Mg2(dobdc) (−41 and −37 vs − 47 kJ mol−1). Hence, Mg2(dobdc) is considered as one of the most prominent CO2 adsorbents, especially in the low-pressure range. Its CO2 and CH4 adsorption capacity at 1 bar and 25 °C is 8.61 and 1.05 mmol g−1, respectively.239,240 The preferential adsorption of CO2 over CH4 is attributed to the ionic character of the MgO bonds.239 Apart from tuning the adsorption capacity, the choice of metal also has an impact on the performance stability. Results of CO2 breakthrough experi-

ments at 70% relative humidity (RH) revealed a poor performance of Mg2(dobdc) as compared to Ni2(dobdc) and Co2(dobdc). The Mg2(dobdc) suffered a sharp reduction in its CO2 adsorption capacity when exposed to humidity. Also, it retained only 16% of the initial CO2 capacity after thermal regeneration, but Ni2(dobdc) and Co2(dobdc) were able to retain ∼60 and ∼85% of their initial capacities, respectively.241 Therefore, it is possible to engineer MOFs with open metal sites by tailoring the metal nodes to meet different requirements, whether it is for efficient adsorption where high CO2 uptake is necessary or for more reliable performances under specific operating conditions. 2.1.2.4. Pre- and Postsynthetic Functionalization. Grafting of functional groups to give a porous surface of high CO2 affinity is an effective approach in enhancing CO2 uptake capacity and CO2/CH4 selectivity. This can be realized by a pre- or postsynthetic functionalization. MOFs that are functionalized by a presynthetic modification usually depend upon engineered ligands as the building blocks to graft desired pendant groups such as −Br, −NH2, −CH3, or other small substituents within the frameworks.223 A significant weakness of this approach is the compatibility issue of the engineered ligands with the typical synthetic methods used for the preparation of MOFs. Postsynthetic functionalization of MOFs can serve as a potential solution to this weakness. To date, a wide variety of functional groups have been introduced using this approach, on the basis that the parent MOFs are sufficiently robust and porous for the precursors to diffuse in. The drawback is the reduction in surface area and sacrificial decrease in adsorption kinetics brought by the introduction of longer and bulkier functional groups in the pore apertures. Q


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Figure 13. (a) Structure of NH2-MIL-53 with O, C, and N atoms in red, gray, and blue. CO2 and CH4 uptake of NH2-MIL-53 at 30 °C, revealing a CO2 hysteresis at high pressure but none for CH4. Reprinted with permission from ref 246. Copyright 2009 Elsevier. Reprinted from ref 242. Copyright 2009 American Chemical Society. (b) CO2 uptakes of Cu-BTTri, en-Cu-BTTri, and mmen-Cu-BTTri. Reprinted with permission from ref 243. Copyright 2009 Royal Society of Chemistry. Reprinted from ref 244. Copyright 2009 American Chemical Society. (c) Structures of Mg2(dobdc) and Mg2(dobpdc) with an increase in pore size from 13 to 21 Å. Diamine functionalization in the pores of mmen-Mg2(dobpdc) affords a higher CO2 uptake at pressure ZIF-82 (−CN, 12.3 Å), ZIF-81 (−Br, 7.4 Å), ZIF-69 (−Cl, 7.8 Å) > ZIF-68 (−C6H6, 10.3 Å), ZIF-79 (−CH3, 7.5 Å) > ZIF-70 (−H, 15.9 Å) (Table 12).269 This order was attributed to the increased CO2 affinity of the functional groups in the sIm linkers and not the pore size of the ZIFs. Moreover, ZIF-78 (−NO2) and ZIF-82 (−CN) with strongly polar groups exhibited relatively higher CO2/CH4 selectivities as compared to other isoreticular ZIFs. The results reinforce the fact that the functional groups (rather than molecular sieving effect) have a predominant influence over the CO2 capture behavior of the isoreticular ZIFs. A similar approach was also used by the same group to prepare an isoreticular series of ZIFs (ZIF-300 to -302) with a CHA topology. The ZIFs showed outstanding water stabilities as evident by the effective CO2/N2 separation under both dry and humid conditions (Figure 15d).270 Other approaches, such as postsynthetic metal ions or ligands exchange, have also yielded ZIFs with zeolitic structures. However, this approach is demonstrated only on a few types of frameworks with incomplete exchange of metal ions or ligands cited as the main reason behind the low success rate.271,272 Despite this drawback, the appealing properties of such ZIFs for CO2 separation continue to drive greater research efforts in novel synthesis and detailed characterization studies. Readers are referred to dedicated reviews264,267,273 to comprehend this subclass of MOFs. As a quick overview, Table 12 summarizes some representative ZIFs with zeolitic structures and the key parameters that are useful for CO2 separation. 2.1.4. Microporous Organic Polymers. Microporous organic polymers (MOPs) are a class of porous materials that are constructed solely by light elements (i.e., H, B, C, N, and O) via strong covalent bonds. MOPs can been classified into hyper-cross-linked polymers (HCPs),279−281 conjugated microporous polymers (CMPs),282,283 porous aromatic frameworks (PAFs),284,285 polymers of intrinsic microporosity (PIMs), 2 8 6 , 2 8 7 and covalent organic frameworks (COFs)288,289 depending on the synthesis approaches and

survived the synthesis in an aqueous medium as well as remained intact in boiling water and methanol for several days. This conclusion was corroborated by TGA and XRD analysis, which suggested uncompromised crystallinity of ZIF-8 even at a temperature as high as 200 °C (Figure 14b).264 Thus, ZIF-8 is a popular choice as a filler for CO2/CH4 composite membranes. To date, more than 100 different types of ZIF structures have been demonstrated.267 Yet, only a few structures possess topologies resembling those of conventional zeolite structures. To address this gap, two synthetic strategies are adopted. The first one uses structure-directing agents as linkers to give geometrically controlled secondary building units that eventually lead to a zeolite topology. For example, Hayashi et al.268 introduced ligands with nitrogen in place of carbon atoms at key positions in the imidazolate linkers. Particularly, ligands with nitrogen atoms placed at position 5 (Figure 15a) were able to exert an electrostatic and dipole−dipole interaction between the HC−N···N−CH pair at positions 5 and 6 of the benzene rings, giving rise to linker−linker interactions that specifically directed the ZIF into a LTA topology (Figure 15b). In the case of 4-azabenzimidazole (Figure 15a), the distance between the two linkers at positions 4 and 7 appeared too far for any favorable interactions, and thus no LTA topology was featured. Essentially, the structuredirecting effect stems from such a favorable interaction, which propagates the ZIF into a LTA structure by forming tilting cubes at the early stage of the crystallization process. The second strategy uses a mixed linker approach. Banerjee et al. prepared an isoreticular series of ZIFs (ZIF-68 to ZIF-70 and ZIF-78 to ZIF-82) using substituted imidazoles (sIm) with functional groups, including −NO2, −CN, −Br, −Cl, −CH3, and −C6H6 (Figure 15c).269 By using equimolar of 2nitroimidazole (nIm) and sIm as a second linker, a GME topology was obtained across the series of ZIFs (Figure 15c). Characterization revealed that the resulting ZIFs showed increasing pore size, ranging from 7.1 to 15.9 Å, with CO2 uptakes (1 bar and 25 °C) in the order of ZIF-78 (−NO2, 7.1 U


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diameter of 6.4, 18.7, and 34.1 Å, respectively (Figure 17a). Controlling the average strut length can also tune the microporosity in amorphous MOPs. For instance, the micropore diameters of the CMPs series (CMP-0 to -5) was systematically enlarged as the strut length increased (Figure 17b).283 However, the enhanced degree of conformational freedom in the extended struts (CMP-3 and -5) allowed for greater space-filling and intermolecular intercalation. Hence, a substantial dip in the overall micropore volume from 0.38 to 0.16 cm3 g−1 was observed as the struts lengthened from CMP0 to -5.282,283 In fact, the microporosity of amorphous CMPs can be finetuned in a continuous fashion using statistical copolymerization of monomers of various strut lengths. Jiang et al.283 prepared a series of CMP copolymers (CPN-1 to CPN-6) by copolymerizing 1,3,5-triethynylbenzene with 1,4-diiodobenzene and/or 4,4′-diiodobiphenyl. A systematic increase in the average micropore size and surface area was observed when a greater portion of the shorter 1,4-diiodobenzene strut was used.283 This method was later expanded to other amorphous frameworks. For example, HCP materials, such as the PP-N-x series, were synthesized by the copolymerization of dichloro-pxylene with triphenylamine. By increasing the mole ratio of dichloro-p-xylene to triphenylamine, the surface area increased from 318 to 1530 m2 g−1.300 At 25 mol % triphenylamine, the ensuing PP-N-25 framework exhibited a high CO2 uptake of 4.60 mmol g−1 at 1 bar and 273 K, owing to the increased surface area and micropore volume (Table 13). The nitrogen atoms in the framework were also able to contribute as a Lewis base to provide a strong affinity toward the CO2 molecules. Similar to MOFs (see section 2.1.2.4), functional groups with high CO2 affinity can be incorporated into MOP networks via presynthetic modifications. To date, several functional groups, such as carboxylic acid (−COOH), amine (−NH2), hydroxyl (−OH), and methyl groups (−CH3), have been incorporated into CMP-1 by reacting 1,3,5-triethynylbenzene with its respective functionalized monomers (Figure 18a).314 Significant increase in the isosteric heats of adsorption and CO2 capture capacities suggested that the polar groups (−COOH, −OH, and −NH2) were effective in enhancing the CO2 affinity of the CMPs but not the nonpolar −CH3 group (Table 13). Functionalization of sulfonate groups into microporous PIs have also been reported using a one-step condensation reaction involving tetrakis(4-aminophenyl)methane (TAPM), binaphthyl dianhydride (BTDA), and sulfonated binaphthyl dianhydride (SBTDA) at different mole ratios.304 The higher CO2 capture capability and CO2/ N2 selectivity of the sulfonated PIs suggested strong CO2 affinity to the −SO3H groups. Postsynthetic modification of MOPs is also a common method to graft CO2-philic functionalities. Among the many MOPs, porous polymer network, PPN-6 (also known as PAF1), is an ideal matrix given its ultrahigh surface area and physicochemical versatility (Figure 18b). The grafting process utilized chlorosulfonic acid to modify PPN-6 into sulfonated PPN-6-SO3H, which could be further neutralized to its lithium salt, PPN-6-SO3Li.240 Due to the strong interaction between CO2 and the sulfonated groups, the CO2 uptake increased from 5.1 to 13.1 and 13.5 wt % for PPN-6, PPN-6-SO3H, and PPN-6-SO3Li, respectively. Both PPN-6-SO3H and PPN-6SO3Li also have small pore sizes (ranging from 5.0 to 10.0 Å), which render them suitable for CO2/CH4 separation. In another study, Lu et al. postsynthetically modified PPN-6 with

the required chemical structures. For HCPs, CMPs, and PAFs, the permanent porosity arises from the numerous linking points between the adjacent building blocks that prevent the polymer chains from collapsing into a dense, nonporous state.290 Conversely, PIMs are porous materials with free volumes derived from space-inefficient packing of highly rigid and contorted polymer chains (see section 1.3.2).291,292 Linear PIMs, such as PIM-1 and -7, are soluble in common organic solvents, and thus processable into defect-free membranes for CO2/CH4 separation.293−296 However, other classes of MOPs lack solvent processability, and so they are generally used as fillers instead. Among the many MOPs, covalent-organic frameworks (COFs) are the only subclass that exhibit a structure of high crystallinity. Similar to MOFs, COFs possess well-defined 2- or 3D ordered porous architectures (Figure 16).222

Figure 16. Building units and structures of (a) 2D COF-1 and (b) 3D COF-2. Adapted with permission from ref 288. Copyright 2005 American Association for the Advancement of Science. Adapted with permission from ref 289. Copyright 2007 American Association for the Advancement of Science.

Typically, MOPs are synthesized by chemical reactions, including metal-catalyzed coupling,283,297,298 oxidative polymerization,299−301 click reaction,302 trimerization,303 and condensation polymerization.304−306 During the covalent bonds formation, effective adjustment of the thermodynamic equilibrium is essential for the formation of highly ordered structures. COFs, on the other hand, are synthesized by several reversible organic reactions, such as the formation of B−O (boronate and boroxine),288,289 CN (imine and hydrazone),306,307 and C−N (triazine and imidization)303,308 bond linkages, to create the necessary building blocks. As covalent bonds are formed, MOPs, as compared to inorganic porous materials (e.g., zeolites and CMSs) and MOFs, possess the advantages of persistent microporosity, strong physicochemical versatility, as well as robust structural and chemical stability, making them highly promising for CO2 capture and CO2/CH4 separation.292,309−311 The porous crystalline structure of COFs can be finely adjusted by changing the strut length of the monomer. This effect is observed through a systematic change in the surface area, pore volume, and/or pore size as the length of the strut varies. Côté et al.312 demonstrated this by extending the spacer length to give COF-6, -8, and -10 with increasing pore V


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Figure 17. (a) Co-condensation of BTBA, BTPA, and BPDA with HHTP to give COF-6, -8, and -10. Adapted from ref 312. Copyright 2007 American Chemical Society. (b) Chemical structure for CMPs with varying strut lengths and the pore size distribution curves derived from N2 sorption isotherms measured at 77 K. A shift to larger average pore size in this series is shown by the dashed arrow. Adapted with permission from ref 313. Copyright 2009 Wiley-VCH. Adapted from ref 283. Copyright 2008 American Chemical Society.

stronger CO2-philic alkylamine groups.315 PPN-6-CH2EDA, PPN-6-CH2TAEA, PPN-6-CH2TETA, and PPN-6-CH2DETA were all obtained by reacting PPN-6-CH2Cl with different alkylamines (Figure 18b). Despite a sharp drop in their surface areas, the polyamine-tethered materials exhibited significant enhancements in their isosteric heats of adsorption and CO2 uptake capacities at low pressures. Particularly, PPN-6CH2DETA showed the highest CO2 uptake capacity (15.8 wt % at 295 K and 1 bar) and CO2/N2 selectivity at 442.315 By and large, MOPs possess appealing parameters and CO2 uptake capacities that are desirable for CO2/CH4 separation (Table 13). Moving forward, greater research efforts should be focused on resolving the challenge of large-scale production and modification of MOPs, evaluating performance stability in mixed-gas separations and creating more robust MOPs frameworks that can prevent the loss of structural integrity when utilized as fillers in composite membranes. 2.1.5. Carbon-Based Particles. Carbon-based particles, herein referred to as carbon bulk materials, exhibit various desirable properties of a good adsorbent which include high porosity, large effective surface area, good thermal and chemical stability, cost-effectiveness, and mature technological

readiness. However, most carbon-based particles, including activated carbons (ACs) and CMSs, lack distinct morphological structures and low-dimensional features that define nanoscale materials such as GFMs and CNTs. In this section, our discussion is focused on ACs and CMSs for CO2 separation. 2D GFMs and 1D CNTs will be discussed later (see sections 2.2.1 and 2.3). 2.1.5.1. Activated Carbons. ACs contain small carbon layer stacks arranged in a poorly crystalline manner.323 The structure of ACs can be described as a network of twisted and defective planes of carbon layers that are cross-linked via aliphatic bridging groups.324 Generally, carbonaceous precursors such as coal, polymer, and biomass are used to obtain ACs through an activation process either by physical or chemical means. Physical activation involves oxidation in air or carbonization under an inert atmosphere at high temperatures while chemical activation uses activating agents like acids, bases, or chloride salts to separate the carbon layers and afford the porous structures.325 Slit-shaped micropores are typically produced with the occasional formation of meso- and macropores (depending on the precursors) to give large pore volume, pore size (∼20 Å) and Brunauer−Emmett−Teller (BET) surface W


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Table 13. Overview of the Key Parameters of Representative MOPs that Show Promise in CO2 Separation MOP

crystallinity


pore size (nm)



ref

CMP-1 CMP-1-COOH CMP-1-NH2 CMP-1-(CH3)2 CMP-1-(OH)2 COF-1 COF-5 COF-6 COF-8 COF-10 COF-102 COF-103 NUS-2 NUS-3 PAF-1 (PPN-6) PIM-1 PPI-1 PPI-2 PPI-3 PP-P PP-PO PP-N-0 PP-N-25 PP-N-50 PP-N-75 PP-N-100 PTNPM PTNPP PTNPM-F PTNPP-F SMPI-0 SMPI-10 SMPI-50 SMPI-100 SNW-1 TpBD TpPa-1

amorphous amorphous amorphous amorphous amorphous crystalline crystalline crystalline crystalline crystalline crystalline crystalline crystalline crystalline amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous amorphous

− − − − − 0.98 0.58 1.10 0.71 0.48 0.43 0.43 − − − − − − − − − − − − − − − − − − − − − − − − −

− − − − − 0.90 2.70 0.60 1.90 3.40 1.20 1.20 0.90 2.00 − − − 5.90 8.00 1.60 1.50 − − − − − − − − − 0.60, 1.40d − − − 0.50 2.40 1.80

1.20 1.10 0.90 0.90 0.90 1.16 0.70 1.93 0.75 0.61 0.77 0.86 0.80 2.50 2.05b 1.20c 1.92 1.56 1.34 3.80b 2.50b 3.90b 4.60b 3.86b 3.13b 2.54b 2.69 1.96 1.60 1.17 1.43 1.87 1.61 1.87 2.05 1.64b 3.48b

0.45 0.30 0.39 0.75 0.71 0.30 1.07 0.32 0.69 1.44 1.55 1.54 0.31 0.43 1.70 0.48 0.04 0.29 1.00 1.11 1.07 2.27 1.05 1.08 0.78 0.56 1.58 0.63 0.06 0.01 0.40 0.12 0.07 0.04 1.54 0.46 0.40

314 314 314 314 314 316 316 316 316 316 316 316 317 317 125 318 319 319 319 298 298 300 300 and 320 300 and 320 300 and 320 300 301 301 301 301 304 304 304 304 321 322 322

a

CO2 feed condition: 1 bar and 298 K. bCO2 feed condition: 1 bar and 273 K. cCO2 feed condition: 1 bar and 308 K. dTwo different pore sizes are present.

areas (as large as 3000 m2 g−1).326 Similar to nonporous carbon blacks, ACs can be functionalized with surface oxides, −NH2, and chemisorbed halogens.325 Besides, they possess a low isosteric heat of adsorption, owing to a low adsorption strength with the adsorbate. For these reasons, ACs are easier to regenerate and deemed highly suitable for CO2 separation. Siriwardane and co-workers327 compared the CO2 adsorption capacities of ACs, molecular sieves 13X and 4A. Results showed that ACs exhibited the highest CO2 adsorption capacity of 8.5 mmol g−1 at 20.7 bar and 25 °C. For biogas upgrading, Castrillon et al.328 recently investigated the potential of ACs impregnated with K2CO3, NaOH, or Fe2O3. ACs impregnated with NaOH were demonstrated to have high CO2 and H2S adsorption capacities of 1.7 and 4.6 mmol g−1, respectively, at 298 K and 1 bar under dry conditions. The mixed-gas selectivity from a binary gas mixture of CO2/CH4 (30:70 vol %) was observed to be 2.4 at 298 K. The results were attributed to the basic pore surfaces and the narrow micropore size distribution of ACs, asserting the possibility of acidic gases removal by ACs during biogas upgrading. On a

different note, nanoporous carbons prepared by pyrolysis of polyfurfuryl alcohol at 550 °C revealed higher isosteric heats of absorption for CO2 as compared to CH4 and N2.329 On the basis of the experimental adsorption isotherms of these three gases, calculated CO2/CH4 (10:90 vol %) and CO2/N2 (10:90 vol %) mixed-gas separation were found to be 8.6 and 31.9, respectively, at 323 K and 10 bar. 2.1.5.2. Carbon Molecular Sieves. CMSs are more suitable for discriminating gas molecules with small disparity in kinetic diameters, owing to their micropores of 5 to 10 Å with a narrower size distribution as compared to ACs (Figure 19a).330 The carbonaceous precursors of CMSs include petroleum pitch,331−333 resin,334,335 cellulosic sources,336−338 and carbon fibers.339−342 By and large, there are two main approaches for synthesizing CMSs. The first approach utilizes a controlled pyrolysis to create micropores in the precursors. However, it is challenging to ensure absolute reproducibility of the pore size and its distribution due to difficulties in replicating the exact synthesis conditions between different batches of CMSs. The second approach exploits a carbon deposition technique to X


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Figure 18. (a) Reaction scheme of functionalized CMPs. Reproduced with permission from ref 314. Copyright 2011 Royal Society of Chemistry. (b) Structure of PPN-6 and its derivatives. Reproduced with permission from ref 292. Copyright 2014 Springer.

Figure 19. (a) Comparison of the pore size and its distribution between ACs and CMSs. Adapted with permission from ref 347. Copyright 1977 Elsevier. (b) Transport mechanism of CO2/CH4 separation, showing adsorption of CO2 through the micropores of CMSs.

In other words, the molecules are required to overcome an energy barrier during penetration. This leads to differences in the adsorption kinetics of the gases, especially when the size of gas molecules is close to that of the micropore entrance.345 The pore size of the CMSs controls the kinetic selectivity, which means that varying the degree of carbon deposition to create appropriately tailored pore apertures is instrumental in ensuring a high CO2/CH4 kinetic selectivity. Also, the adsorption kinetics involves an intricate interplay between the size, shape, electronic structure of the gas molecules and

heterogeneously precipitate carbon onto another microporous carbon substrate such as ACs to create a pore blocking effect.343 To this end, several carbon deposition agents which include acetylene, methane, cyclohexane, and benzene have been employed.344 The key advantage of this approach lies in its ability to tailor the pore dimension close to the size of the desired adsorbate such as N2 (0.364 nm) and O2 (0.346 nm) in air separation. It is generally accepted that the diffusing gas molecules experience net repulsions when accessing the pores of CMSs. Y


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Figure 20. Liquid crystal templating (LCT) process. Reprinted with permission from ref 354. Copyright 2012 Royal Society of Chemistry.

from 2 to 50 nm. At first glance, these pore diameters appear too large and out of the range of interest for CO2/CH4 separation. However, mesoporous materials also possess huge pore volumes in addition to their tunable mesopores, which can be valuable for improving the limited accessibility of gas molecules as observed in zeolites and MOFs with smaller micropore channels.352 Currently, mesoporous silicates, such as MCM-41, MCM-48, SBA-15, and COK-12, dominate because their mesoporous structures are relatively easier to control by a surfactant templating method in contrast to the nonsilicate counterparts.353 On top of that, other advantages of mesoporous silicates include large surface areas (above 700 m2 g−1), narrow pore size distributions, uniform and tunable pore sizes, as well as low mass densities.354 More importantly, pore engineering can potentially realize 3D pore systems with enhanced CO2 mass diffusion355 and functionalized mesopores with increased CO2 affinity.354 Some of the earliest work on ordered mesoporous molecular sieves was pioneered by researchers from Mobil.356 The family of mesoporous materials known as M41S consists of MCM-41 and -48 (MCM: Mobil Composition of Matter), which were synthesized using a cationic surfactant and a solubilized anionic silica source via a liquid crystal templating (LCT) method (Figure 20).357,358 The MCM-41 structure possesses a 2D hexagonal system with a p6mm space symmetry and unidirectional mesopores in the range from 15 to 100 Å (Figure 20). The MCM-48 structure, on the other hand, exhibits a cubic system with a Ia3d space symmetry and a bicontinuous pore structure with mesopores in the range from 15 to 30 Å (Figure 21a). Heuristically, the synthesis of both MCM-41 and -48 depends on the ratio of the silica source to surfactant.359 A

the interaction with the pores of the CMSs. Particularly, the CO2 molecule has a linear shape and smaller kinetic diameter as compared to the larger tetrahedral-shaped CH4 molecule (Table 2). It is therefore expected that CO2 molecules diffuse through the pores along its minimum dimension (Figure 19b).343 The favorable transport is also driven by a more significant intermolecular interactions given the strongly quadrupolar nature of CO2 molecules (Table 2). As such, the adsorption kinetics of CO2 is faster than CH4, suggesting a strong relevance of CMSs for CO2/CH4 separation.346 Similar to ACs, the CO2 adsorption capacity of CMSs is also often benchmarked against conventional porous materials such as zeolite 5A and 13X. Apart from displaying strong thermal and chemical stabilities,348 CMSs are less affected by the presence of water due to its hydrophobic surface, making it an effective filler for addressing biogas upgrading under realistic operating conditions.327,349,350 Compared to zeolites with typical surface areas of 500 to 1000 m2 g−1, CMSs synthesized from petroleum pitch have higher porosities and BET surface areas that reach as high as 3100 m2 g−1.331 Despite that, the assynthesized CMSs demonstrated a lower CO2 uptake than zeolites due to a relatively dominant adsorbent−adsorbate interaction by zeolites at the low CO2 partial pressure regime. At a higher pressure of 1 bar and 273 K, the CO2 adsorption capacity of CMSs, however, surpassed zeolite 5A (∼4.1 mmol g−1) and 13X (∼5.0 mmol g−1) with a reported value close to 9.0 mmol g−1. Correspondingly, the CMSs exhibit effective molecular sieving effect for the CO2/CH4 gas pair. Characterization unveiled a micropore size entrance of less than 0.4 nm, which afforded a high CO2/CH4 selectivity of 14 measured at 298 K. The same group later extended the study to a wider range of pressure from atmospheric pressure to 50 bar. At a low pressure of 1 bar, the as-synthesized CMSs exhibited a CO2 uptake of ∼4.5 mmol g−1 at 298 K. This value was 4- to 6fold greater than that of MOF-200 and -177. At a high pressure of 50 bar and considering volumetric adsorption, the CMSs displayed a remarkable CO2 adsorption of nearly 500 cm3 (STP) per cm3 at 298 K.351 Contrastingly, adsorptions by MOF-177 and MIL-101 were only 320 and 390 cm3 (STP) per cm3, respectively. The higher volumetric adsorption was attributed to the smaller packing density of CMSs as compared to the higher crystal density of MOFs, which is a clear advantage of CMSs at the high-pressure regime. 2.1.6. Mesoporous Materials. The IUPAC defines mesoporous materials as particles with pore diameters ranging

Figure 21. Structures of the various mesoporous silicates. (a) MCM48. Reprinted with permission from ref 383. Copyright 2006 WileyVCH. (b) SBA-15. Reprinted with permission from ref 384. Copyright 2009 Royal Society of Chemistry. (c) COK-12. Image courtesy of Prof. C. Liang. Z


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Figure 22. Chemical structures of (a) pristine graphene, (b) graphene oxide (GO), and (c) reduced GO (rGO).

interconnected cylindrical pores within the primary mesopores370,371 as substantiated by its wide pore size distribution.372 The cross-channel connectivity is the key to a fast CO2 transport.373−378 However, due to the larger pore channels, SBA-15 possesses a lower degree of capillary condensation and generally shows weaker CO2 adsorptions before and after amine-functionalization. In particular, as compared to MCM41 with CO2 adsorption of 0.67 mmol g−1 before and 1.16− 2.20 mmol g−1 after functionalization, SBA-15 only displayed CO2 adsorption of 0.04−0.59 mmol g−1 before and 0.52−1.40 mmol g−1 after functionalization.379 Another noteworthy mesoporous silicate, COK-12, has ordered 2D p6mm hexagonal structure with straight pore channels similar to MCM-41 and SBA-15 (Figure 21c). Most studies utilize COK12 as fillers for composite membranes for gas separation,380 which will be discussed in section 4.1.6. These studies are often motivated by the simple synthesis of COK-12 at room temperature and under quasi neutral pH value.381,382 A major drawback of mesoporous silicates is their low hydrothermal stabilities. The siloxane bonds are prone to hydrolysis when exposed to moisture at elevated temperatures, leading to weak structural stability against moisture and easy collapse of the pore walls.379 This is exemplified by Cassiers et al., who showed a 15 to 20% decrease in the mesopore volumes of SBA-15 and MCM-48 along with a 40 to 50% higher decline for MCM-41 after a mild hydrothermal treatment at 400 °C and 1 bar with 25% steam.385 Present strategies to improve the hydrothermal stability of mesoporous silicates involve thickening of the pore walls and increasing the hydrophobicity of the pore through silylation and alumination to reduce exposure of the silanol groups to moisture.386,387 However, such modifications are usually accompanied by alterations to the textural properties, and more work is needed to establish the impacts of such modifications on the CO2 adsorption performance.388

ratio of less than 1 usually favors the formation of the MCM41 structure. The LCT process typically proceeds with a selfassembly of the cationic surfactant into micelles templates, which then serve as structure-directing agents to afford the MCM-41 crystallites. However, the exact mechanism remains controversial and requires additional investigation (see refs 352 and 360 for the arguments). Upon successful crystallization of the hexagonal array, calcination or extraction is generally employed to remove the excess surfactants and open the pores, leaving behind hollow 1D channels of inorganic materials (Figure 20). The former approach is normally preferred given its facile implementation.361 MCM-41 is highly versatile in terms of its compositional makeup and physicochemical properties. For instance, aluminophosphates or other transition metal oxides, including that of Fe, Sb, Zr, Ti, V, Mn, and W, can be used for the effective synthesis of MCM-41.188,362 The synthesis conditions, such as concentration, temperature, and pH value, play significant roles in ensuring the stability of the supramolecular structure. These parameters are equally important for pore size control353 alongside pore tailoring strategies by taking advantage of cationic surfactants with different alkyl chain lengths,358,363,364 micellar swelling365 or the addition of auxiliary hydrocarbons like 1,3,5-trimethylbenzene.366 The hydrophilicity of MCM-41 is also tunable by varying the Si/ Al ratio to a certain extent. As a result, MCM-41 demonstrates favorable physicochemical properties for CO2 adsorption.367,368 The major issue, however, is the 1D pore structure of MCM-41, which makes the channels more susceptible to pore blockage and, in turn, reduces interaction with the adsorbates. On this account, MCM-48 appears more advantageous owing to its 3D pore system with facilitated mass diffusions.359 The mesopores of both MCM-41 and -48 can also be postsynthetically functionalized to enhance their CO2 uptakes. As reported by Mello et al.,368 amine-grafted MCM41 exhibited substantially higher CO2 uptake of 0.70 mmol g−1 at 0.1 bar and 30 °C as compared to nongrafted MCM-41 (0.12 mmol g−1) under the same conditions. Similar behavior was also observed by Huang et al.369 on surface-functionalized MCM-48 with around 7-fold increase in CO2 uptake at 25 °C and 1 bar. SBA-15 is another mesoporous material with the capacity to adsorb CO2. The synthesis of SBA-15 employs amphiphilic triblock copolymers instead of cationic surfactants as the template to yield the mesoporous structure. Interestingly, the resulting structure of SBA-15 is comparable to MCM-41. It has a 2D hexagonal system with p6mm space symmetry and cylindrical pores that are tunable up to 300 Å (Figure 21b).366 Yet, SBA-15 is distinct in the sense that it possesses

2.2. Two-Dimensional Materials

2.2.1. Graphene-Family Materials. Graphene refers to 2D nanomaterials comprising a single layer (1-atom-thick) of sp2-hybridized carbon atoms arranged in a hexagonal honeycomb-like lattice (Figure 22a). Given that the resistance of a molecule across the membrane is inversely proportional to its permeation path (i.e., membrane thickness), graphene is conceptually a membrane with practically the lowest possible resistance.389 Graphene also possesses remarkable mechanical strength, which can be exploited through graphene oxide (GO, an oxidized form of graphene) and reduced graphene oxide (rGO) (Figure 22b,c).390 Both derivatives of graphene are highly tunable and processable in solutions, and so they have attracted much attention from the membrane community.391,392 AA


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Figure 23. Schematic illustration showing (a) pore generation to give size-selective single-layer graphene. Reprinted with permission from ref 401. Copyright 2013 American Association for the Advancement of Science. (b) He atom deflecting from a graphene surface with a kinetic energy of 18.6 eV; (c) models of defective graphene with different vacancies. Reprinted with permission from ref 393. Copyright 2014 Wiley Periodicals, Inc. (d) Ion bombardment and (e) Au nanoparticle deposition on a graphene sheet. Reprinted from ref 397. Copyright 2014 American Chemical Society. Reprinted from ref 400. Copyright 2015 American Chemical Society.

Figure 24. Schematic illustration of (a) an aligned multilayered laminate demonstrating the highly tortuous transport pathway as a result of the impermeability of GFMs. Reprinted with permission from ref 393. Copyright 2014 Wiley Periodicals, Inc. (b) An increased interlayer spacing between GO sheets when intercalated by PEI. Reprinted from ref 407. Copyright 2011 American Chemical Society. (c) Interfacial void between graphene and the PDMS matrix in the composite membrane. Reprinted from ref 117. Copyright 2015 American Chemical Society.

2.2.1.1. Single-Layer Graphene. Before utilizing a singlelayer graphene as membrane for gas separation, it is necessary to first generate nanopores on the surface of the graphene (Figure 23a). This is because the electron density of the aromatic rings on a defect-free single-crystalline graphene is so large that it can repel atoms and molecules as demonstrated through MD simulations (Figure 23b).393,394 Even for defective graphene sheets, which possess Stone−Wales (SW) and di- to decavacancy defects (Figure 23c), the energy barriers remain too high for an He atom to pass through.395 Therefore, to allow molecules to permeate a single-layer graphene, subnanometer pores are essential on the graphene surface to lower the energy barrier and facilitate molecular transport. As a proof-of-concept, Jiang et al. revealed in their MD simulations that all-hydrogen passivated pore as small as 0.25 nm can allow H2 permeation with an easily surmountable energy barrier of only 0.22 eV and yet provides a high energy barrier of 1.6 eV for the bigger CH4 molecules. Resultantly, an extremely high H2/CH4 selectivity on the order of 1023 was observed. For a nitrogen-functionalized pore at 0.3 nm, the selectivity decreased slightly but remained in the order of 108 with a high H2 permeance.396 Precise pore size control can therefore technically tune the selectivity of the single-layer

graphene membrane while achieving high gas permeability. To this end, nanopores on single-layer graphene have been experimentally realized by etching techniques assisted via ion bombardment, Au nanoparticle deposition, O2 plasma, and UV-induced oxidative treatment (Figure 23d,e).397−400 2.2.1.2. Multilayer GFMs. Beyond single-layer graphene, multilayered GFMs are also used as laminates or fillers in separation membranes.392 The key motivation is to combine the benefits of 2D structures and virtual in-plane impermeability of GFMs to realize a highly tortuous pathway for the diffusion of gas molecules through the interlayer spacing of the laminates or around the fillers (Figure 24a).393 In an ideal situation where the GFMs sheets are orientated perpendicular to the flow direction of the molecules, the diffusion pathway taken by the molecules can literally extend over 1000 times the thickness of the laminate.402 Indeed, many studies focusing on GFMs/polymer composites have demonstrated lower gas permeances as compared to their purely polymeric counterparts, suggesting a high barrier effect of the GFMs.403 The lowered gas permeance is often accompanied by an increase in the membrane selectivity.404 For instance, Yang et al. employed a layer-by-layer (LbL) assembly method to fabricate a 10-bilayer GO/polyethylenimine laminate (∼91 nm in AB


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thickness) with an order of magnitude lower CO2 transmission rate but a much improved H2/CO2 selectivity that reached above 383. 405 They ascribed the observation to the impermeability of the GO sheets, which resulted in a highly extended diffusion pathway of CO2 through the brick-andmortar architecture. The polyethylenimine (PEI), on the other hand, served as a spacer to increase the interlayer spacing between the GO sheets and facilitated a relatively lower H2 transport resistance (Figure 24b). Similarly, Huang et al. fabricated a GO/poly(lactic acid) (PLA) composite membrane with reduced O2 and CO2 permeability coefficients. The results were in agreement with a modified Nielsen’s model, suggesting that GFMs continued to exert good barrier effect even when incorporated in a polymer matrix.406 Some studies, however, gave contradictory results showing increased gas permeability, which may or may not be accompanied by a decline in membrane selectivity. Usually, such an increase in gas permeability is attributed to the presence of interfacial voids between the polymer matrix and GFMs filler. The voids create greater free volumes in the polymer matrix and offer supplementary pathways of lower transport resistance for the gas molecules (Figure 24c).117 2.2.1.3. Functionalized GFMs. Chemical functionalization is certainly one of the most promising strategies for enhancing the physicochemical properties of GFMs for CO2 separation membranes. Functionalization including modifying the graphene surfaces and designing highly porous 3D architectures from nonporous 2D GFMs are generally targeted at improving the CO2 affinity, increasing the surface area and optimizing the accessibility of the pores of the GFMs.408 For example, Srinivas et al. synthesized GO derived carbons (GODCs) by KOH activation of thermally exfoliated GO and solvothermally exfoliated rGO sheets. The GODCs demonstrated a range of BET surface areas (up to 1900 m2 g−1), pore volumes (up to 1.65 cm3 g−1), and pore sizes and distributions, which can be tailored by varying the GO/KOH concentrations and activation temperatures. More importantly, the micro- and mesoporosity of GODCs were tunable, making them favorable for CO2 adsorptions.409 GODCs can also be doped with polypyrrole and polythiophene to give N- and S-doped GODCs.410,411 The N-doped GODCs exhibited a highly selective adsorption of CO2 over N2 (4.3 vs 0.27 mmol g−1 at 298 K and 1 bar), owing to the interactions between CO2 molecules and the N-containing functional groups (Table 14).410 Similarly, the S-doped GODCs reported a high CO2 adsorption (4.5 mmol g−1 at 298 K and 1 bar) over N2, CH4, and H2 gases, and showed stable adsorption of 4.0 mmol g−1 upon recycling. On the whole, the enhanced CO2 uptake of Sdoped GODCs was attributed to the high surface area (1396 m2 g−1), optimized S content and tunable microporosity. These key parameters are overviewed together with CO2 uptakes of other representative functionalized GFMs (Table 14). 2.2.1.4. Templated GFMs. A templated approach to synthesize GFMs is also promising in obtaining high surface area and microporosity. Ning et al.420 utilized a chemical vapor deposition (CVD) method and a porous MgO template to prepare nanomesh graphene (NMG) with a BET surface area of 2038 m2 g−1, nanopore of ∼1 nm, and pore volume of 2.35 cm3 g−1. Due to these attractive parameters, high CH4 and CO2 adsorptions of 14.5 (90 bar) and 36.5 mmol g−1 (31 bar), respectively, were obtained at 273 K.420 Recently, the choice of template has been widened to include zeolites and MOFs. The

Table 14. Overview of the Key Parameters of Representative Functionalized GFMs that Show Promise in CO2 Separation GFM degassed GO exfoliated GO exfoliated GO exfoliated nanoplates GO frameworks GO-like foams GO-Mn3O4 (1:4) GO-PEI hydrogel hydrogen exfoliated GO KOH-activated GO KOH-activated Ndoped GO thermal exfoliated GO hydrothermal rGO KOH-activated and S-doped rGO


pore volume (cm3 g−1)

ref

477 31 701 480

0.70b 1.70b 39.00c 56.00c

1.04 0.04 − 1.73

412 413 414 414

470 510 541 253 443

2.70d 2.00e 2.50 2.50b 21.60f

− 0.89 0.31 0.70 −

415 416 417 413 418

1894 1360

2.00 4.00g

1.60 0.59

409 410

925

8.60h

1.18

419

876 1396

1.84b 4.50

− 0.82

413 411

BET surface area (m2 g−1)

a CO2 feed condition: 1 bar and 298 K. bCO2 feed condition: 1 bar and 273 K. cCO2 feed condition: 30 bar and 298 K. dCO2 feed condition: 4 bar and 300 K. eCO2 feed condition: 10 bar and 298 K. f CO2 feed condition: 11 bar and 298 K. gCO2 feed condition: 1 bar and 300 K. hCO2 feed condition: 1 bar and 195 K.

porous carbon structures are usually achieved by carbonizing either the organic linkers in the MOFs or the small organic moieties from a second precursor to yield graphenic fragments with hierarchical micro- and mesopores.421−426 Notably, the intrinsic porosity of the templates is mostly kept intact after carbonization (Figure 25).422,424 As a result, carbon structures possess large surface areas (up to 3300 m2 g−1) and total pore volumes (up to 5.53 cm3 g−1) that are favorable toward CO2 uptake (Table 15).421,424 2.2.2. Layered (Lamellar) Silicates. Layered silicates or clays are naturally occurring hydrated layered aluminosilicate minerals that are present in the form of composite layers or sheets. They are generally constructed based on sheets of tetrahedrally coordinated Si along with edge-shaped octahedral sheets made of Al. The thickness is ∼1 nm, while the lateral dimension ranges from 30 nm to several microns, depending on the type of silicate.428 Hence, layered silicates generally possess high aspect ratios. It is possible to substitute the cations of layered silicates to give negative charges in the overall frameworks. Such a process is termed isomorphic substitution, where partial replacement of cations occurs in both the tetrahedral and octahedral layers. In most cases, Al3+ ions are used to replace Si4+ in the tetrahedral layers, whereas Li+, Mg2+, or Fe2+ ions are used for the octahedral layers. On the basis of the valency of the replaced cation in the octahedral layers, layered silicates can be defined as dioctahedral (twothirds of the sites are occupied by trivalent ions) or trioctahedral (all sites are occupied by divalent ions). There are two main categories of layered silicates: (1) twolayered silicates comprising an octahedral layer linked to a tetrahedral sheet (1:1 type) and (2) three-layered silicates that consist of an octahedral layer sandwiched between two tetrahedral sheets (2:1 type). In both types of silicates, the layers are held by weak van der Waals forces or electrostatic forces of attraction, and the negative charge of the framework AC


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Figure 25. Schematic illustration showing intact frameworks of templated GFMs by (a) porous MgO. Reprinted with permission from ref 420. Copyright 2012 Royal Society of Chemistry. (b) ZIF-8 on PS nanoparticles. Reprinted with permission from ref 427. Copyright 2014 Royal Society of Chemistry. (c) MOF-5. Reprinted from ref 424. Creative Commons license CC BY 3.0.

Beyond aluminosilicate, titanosilicate is another type of silicate which can assume a lamellar configuration. Classical examples include Aveiro-Manchester (AM), JDF-L1, and ETS. The layered forms of AM and JDF-L1 demonstrate considerably high surface areas, although the 2D structure of ETS has yet to be reported (see section 2.1.1.3).433 We contend that future work should focus on investigating the CO2 uptakes of these materials. This is because 2D titanosilicates often exhibit pore sizes that are comparable to the kinetic diameters of CO2 and CH4 but few studies exploit layered titanosilicates as adsorbents or fillers. An example is AMH-3. It is a 3D microporous layered silicate with 8membered ring apertures and crystallographic pore size of 0.34 nm (rings of eight Si−O−Si units).434,435 However, the use of AMH-3 has been challenged by issues such as swelling of AMH-3 and poor preservation of the pore structures and aspect ratios during exfoliation of the AMH-3 layers.435 Other examples also include swollen materials of JDF-L1 and AM-4, which possess lamellar structures with appealing interlayer spacing for CO2 separation.436,437

Table 15. Overview of the Key Parameters of Representative Templated GFMs that Show Promise in CO2 Separation template

BET surface area (m2 g−1)



ref

MgO MOF-5 PS/ZIF-8 ZIF-69

2038 2734 1724 2264

36.50b 27.40c 4.33 4.76d

1.86 5.53 0.63 1.18

420 424 427 426

a

CO2 feed condition: 1 bar and 298 K. bCO2 feed condition: 31 bar and 274 K. cCO2 feed condition: 30 bar and 300 K. dCO2 feed condition: 1 bar and 273 K.

is compensated largely by exchangeable Group IA or Group IIA cations typically found in the interlayer spacing between the layered silicates (see section 2.1.1).429 Hence, the interlayer spacing (a.k.a. gallery) is usually susceptible to intercalation by water molecules, resulting in swelling or even exfoliation of the stacked layers. Swelling of layered silicates is strongly dependent on the type of exchangeable cations.430 For instance, swelling is prominent for small monovalent cations such as Na+ and Li+ but appears otherwise for polyvalent cations. This is because polyvalent cations like Ca2+ can induce chemical cross-linking to prevent the silicate layers from delaminating.430 More recently, 2D layered aluminosilicates have garnered much attention. The first series of layered aluminosilicate is MCM-22, which is a 10-membered ring aluminosilicate comprising 3D pores after removal of the structure directing templates. Modifications to MCM-22 afford either an exfoliated structure, ITQ-2, or a mesoporous structure, MCM-36 (Figure 26a), which CO2 uptakes are strongly affected by factors such as the Si/Al ratio (see section 2.1.1). Layered aluminosilicate based on the MFI framework was also synthesized by a bifunctional quaternary ammonium-type surfactant that possesses a long chain alkyl group, C22. The diammonium head was exploited to direct the formation of a crystalline microporous MFI zeolite, whereas the C22 hydrophobic group helped to suppress the formation of a 3D structure. Consequently, a lamellar surface was formed (Figure 26b) with a resulting pore structure resembling that of a conventional MFI (0.55 nm).431 The extended long C22 chain enlarged the interlayer spacing and gave rise to an intrinsically swollen structure, making the materials highly relevant as fillers for composite membranes.

2.3. One-Dimensional Materials: Carbon Nanotubes

CNTs are nanosized, needle-like hollow cylinders conceptually made up of rolled-up graphene sheets. A CNT with only one graphitic shell is called single-walled carbon nanotube (SWCNT), while one with multiple graphitic shells is called multi-walled carbon nanotube (MWCNT) (Figure 27). CNTs possess many characteristics, one of which is the exceedingly high aspect ratios.438,439 This means CNTs possess the potential to be stretched up to several millimeters in length while retaining only a few nanometers in diameter, making them ideal as nanochannels for gas adsorption and separation. Also, the selectivity and flux for light gas transport in CNTs are a few times higher than other materials with comparable pore sizes.440,441 Such exceptional performances stem from the fast molecular diffusion through the low-friction inner graphitic walls of the CNTs442 and the capacity to offer selective transport through the well-defined diameter of the CNTs.443 2.3.1. Single-Walled Carbon Nanotubes. In principle, SWCNT is constructed from a rolled-up graphene sheet. Given a particular rolled-up direction, a chiral index (n,m) can be used to label a SWCNT and define its symmetry and diameter (Figure 28a). As a result, there are many possible types of SWCNTs. Quiñonero and co-workers embarked on a DFT simulation to comprehensively study the CO2 adsorption of SWCNTs of different diameters, symmetries, and binding AD


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Figure 26. (a) Reaction schemes of MCM-22, MCM-36 and ITQ-2. Reprinted with permission from ref 432. Copyright 2012 Elsevier. (b) Formation of the layered MFI via a surfactant-assisted approach. Reprinted with permission from ref 160. Copyright 2009 Nature Publishing Group.

interior is energetically more favorable than interstitial adsorption for a SWCNT bundle (Figure 28b). In another study, Cinke et al. experimentally evaluated the adsorption of purified HiPCo SWCNTs and showed a CO2 uptake of 5.2 mmol g−1 at 273 K and 1 bar (Table 16). Comparatively, the uptake by SWCNTs was nearly twice that of the ACs despite having merely a 25% higher surface area. The CO2 heat of adsorption was also computed at 0.024 eV. On this basis, it was established that CO 2 molecules were most favorably physisorbed on the side of the SWNCTs, which was contradictory to the simulation results by Quiñonero and coworkers.447 Zhao and co-workers also carried out simulation studies using a combination of techniques that include DFT and local density approximation (LDA) to investigate the adsorption of H2, N2, and CH4 gases by SWCNTs. Analogous to the CO2 adsorption, these gases appeared weakly bound to SWCNTs with physisorption identified as the primary interaction between the nanotube and gas molecule. The optimal binding energy of H2 gas on a (5,5) SWCNT was computed at 0.84 eV, while for N2 and CH4 gases, the binding energies were found to be close to that of CO2 gas at 0.123 and 0.122 eV, respectively.448 They also reported increased CO2 interaction when the nanotubes are in a bundle, pointing to the fact that there is stronger gas adsorption at the interstitial and groove site of a SWCNT bundle than on the surface of an individual SWCNT (Figure 28b). Similar to the results obtained by Cinke et al.447 but contradictory to that by Quiñonero and coworkers,446 the gas adsorption of H2, N2, and CH4 was

Figure 27. Schematic illustrations of a (a) SWCNT and (b) MWCNT showing the number of graphitic walls. HR-TEM micrographs showing (c) SWCNTs in a bundle. Reprinted with permission from ref 444. Copyright 1997 Nature Publishing Group. (d) Two MWCNTs with different number of graphitic walls. Reprinted with permission from ref 445. Copyright 1992 Nature Publishing Group.

sites.446 Results showed that both (9,0) and (5,5) SWCNTs were good candidates for CO2 adsorption due to favorable adsorption thermodynamics as compared to the other studied SWCNTs. The computed CO2 binding energies inside and outside the SWCNTs were also optimized at ∼0.55 and ∼0.12 eV, respectively, suggesting that the CO2 adsorption in the AE


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Figure 28. Schematic illustrations showing (a) different rolled-up directions of a graphene sheet to obtain SWCNT of different diameters and symmetries defined by a chiral index of (n,m). Reprinted with permission from ref 449. Copyright 2015 Elsevier. (b) SWCNT bundle with different CO2 binding sites, namely, (1) surface, (2) pore, (3) groove, and (4) interstitial. Adapted with permission from ref 450. Copyright 2009 The American Physical Society.

co-workers also employed atomistic simulations to study the CO2 transport diffusion in the inner channel of SWCNTs of different diameters. Their results showed that the diffusivity of CO2 was extremely rapid and independent of the applied pressures.452 However, their proposed mechanism of surface or configurational diffusion was contrary to the Knudsen-like diffusion as verified by other experimental studies.452−454 2.3.2. Multi-Walled Carbon Nanotubes. As compared to SWCNTs, MWCNTs offer the competitive advantage of costeffectiveness, more accessible channels given the larger inner diameters of MWCNTs and most importantly, the faculty to tune the physicochemical properties for CO/CH4 separation. To highlight these advantages, Rahimi et al. studied a doublewalled CNT (DWCNT) with an inner diameter of ∼5 nm using Grand-canonical Monte Carlo simulations. They observed a majority of the inner tube volume being accessible to CO2 molecules. Also, the inner and outer walls of the DWCNT exhibited CO2 adsorptions that were mutually exclusive when simulated at 303 K under an applied pressure of less than 40 bar. The study was also extended to a vertically aligned DWCNT forest where the interstitial distance between the DWCNTs was found to have a dramatic impact on the CO2 uptake and its mechanism.455 To capitalize on the tunable physicochemical properties of MWCNTs, chemical functionalization such as oxidation or modification with grafting agents is exploited ubiquitously. Such functionalization is harsh and may inflict defects on the outer walls of the nanotubes.456 For this reason, MWCNTs have an edge since SWCNTs contain only one graphitic wall. To date, several grafting agents have been utilized, including urea, aqueous ammonia, PEI, monoethanolamine (MEA), (3aminopropyl)triethoxylsilane (APTES), and N-[3trimethoxysilyl)propyl]ethylenediamine (TMSPEDA).457−460 The grafting agents are all amine-rich, and when attached to the MWCNTs, create new binding sites with higher CO2 affinity.458 As a result, the CO2 uptake of the modified MWCNTs are generally better than the pristine ones (Table 16). However, the CO2 uptakes tend to differ, even among MWCNTs with the same chemical modification. For example, APTES-modified MWCNTs exhibit CO2 uptakes that fluctuate between 0.9 and 1.7 mmol g−1 (Table 16). The different diameters of MWCNTs, which result in different surface areas, pore volumes, and grafting loadings, as well as the different evaluation conditions employed for each study, are likely to contribute toward this fluctuation. Hence, it is important to

Table 16. Overview of the Key Parameters of Representative CNTs that Show Promise in CO2 Separation

CNT SWCNTs (HiPCo) DWCNTs MWCNTs

chemical state as-grown (raw) purified as-grown APTES-modified APTES-modified APTES-modified EDA-modified MEA-modified NH3(aq)-treated PEI-modified as-grown urea-treated

BET surface area (m2 g−1) − 1617.0 720.0 97.4 128.1 198.0 − − − − 394.0 152.8



ref

b

−

447

b

− − 0.75 − 0.63 − − − − 0.91 0.30

447 455 459 460 457 459 457 457 459 457 458

2.2

4.0 7.2c 1.3e 1.7f 0.9 1.0e 0.6 0.7 0.7e 0.5 1.5d

a

CO2 feed condition: 1 bar and 298 K. bCO2 feed condition: 1 bar and 308 K. cCO2 feed condition: 38 bar and 303 K. dCO2 feed condition: 1 bar and 343 K. eCO2 feed condition: 1 bar and 293 K. f CO2 feed condition: 1 bar and 333 K.

energetically less favorable in the pores (inside) of individual SWCNT than at the interstitial sites (outside) of a SWCNT bundle. Moreover, no clear correlation between the adsorption capacity of the SWCNT and its diameter and symmetry was found.448 Fundamentally, these inconsistencies denote the need to critically evaluate the simulation methods as well as the parameters to establish the credibility of the computational results. The CO2 diffusion from the exterior to interior of a (5,5) SWCNT is computed to be thermodynamically feasible, showing clear evidence of a preferential CO2 adsorption into the inner channel of SWCNT.446 Huang et al. demonstrated this by performing a grand-canonical Monte Carlo simulation using an equimolar mixture of CO2/CH4 gas. A much higher CO2 than CH4 adsorption in the inner channel of a SWCNT was observed. The CO2 uptake varied from 4.0 to 9.0 mmol g−1 depending on the diameter of the SWCNT, while the CH4 uptake remained modestly low at below 1.0 mmol g−1.451 Driven by this preferential CO2 adsorption, the CO2/CH4 selectivity was generally above 8 with the highest value at 11.2 using a (6,6) SWCNT at 343 K and 10 bar.451 Skoulidas and AF


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theless, poor adhesion at the polymer−filler interface is a prevalent issue of concern, especially at higher filler loadings.475 This brings us to the next section where we discuss the different types of nonideal interfacial morphologies which can occur during composite membrane fabrication. The strategies to overcome nonidealities in composite membranes using nonporous fillers will also be unveiled in section 4.4. As a conclusion to this part, we have extensively reviewed a list of conventional and emerging materials with potential as effective fillers for composite membranes for biogas upgrading. The inherent attributes of each type of filler have been covered, including surface area, pore size and distribution, dimensional and morphological structure, chemical and thermal stability, as well as the chemical properties that enable functionalization for achieving enhanced CO2 affinity and adsorption. Table 17 organizes all these essential attributes according to the materials which demonstrated them, alongside their respective limitations and improvement strategies. The table also serves as a useful reference for readers wanting to gain an overview and rapid familiarity with the topic.

consider these differences along with various intrinsic parameters of the CNTs (Table 16) to provide a critical evaluation of the data and an accurate measure of the impact of chemical functionalization on the CO2 adsorption of the CNTs. 2.4. Nonporous Materials

Contrary to the conventional filler materials as described in section 2.1, nonporous materials have no intrinsic pores for gas transport facilitation and molecular sieving effect. Surprisingly, promising membrane separation performances are observed when they are used as fillers in composite membranes. One reason could be due to the favorable surface properties of the nonporous materials where CO2 affinity gets enhanced. More significantly, nonporous materials can disrupt the polymer chain packing and alter the free volume element of the matrix to create a profound impact on the gas transport properties of the composite membranes. A more detailed explanation of this impact will be provided in section 4.4. Here, the focus is on the properties of the materials, in particular, fumed silica (a.k.a. fumed SiO2 or nonporous silica) and other metal oxide nanoparticles of interest such as magnesium oxide (MgO) and TiO2. Fumed silica is composed of amorphous silica particles fused into a 3D structure by flame hydrolysis of silicon tetrachloride (SiCl4).461 During the synthesis, the primary particles of the fumed silica (5 to 50 nm) first merge into chainlike branched aggregates (100 to 500 nm), which then cluster to give micronsized agglomerates.462 Although fumed silica produced by this process has a highly agglomerated structure, a wide range of surface areas from 50 to 500 m2 g−1 remains accessible to gas molecules.463 Furthermore, hydrophilic silanol groups (Si− OH) are readily available on the surface of fumed silica to interact with polar molecules such as CO2.464 These groups can also be easily modified using silane chemistry to give nanospacers around the silica core to enhance CO2 transport.465 Coupled with its cost-effectiveness, fumed silica is therefore unanimously chosen as fillers for composite membranes. Nonporous metal oxide nanoparticles such as MgO and TiO2 also have comparable advantages. Magnesia or MgO is a crystalline material,466 which is synthesized via calcination or thermal decomposition of magnesium hydroxide, Mg(OH)2, or magnesium carbonate, MgCO3.467 Similar to fumed silica, MgO possesses basic surface sites that present strong affinity and interaction toward polar gas species such as CO2.468,469 Compared to other metal oxides with similar chemical properties (e.g., calcium oxide, CaO), MgO exhibits a lower energy penalty for regeneration.470 However, MgO agglomerates easily and is hygroscopic, making it unstable and susceptible to water adsorption when exposed to moisture. In addition, MgO shows a poorer CO2 uptake as compared to conventional fillers like zeolites. For example, zeolite 5A showed a CO2 uptake of 4.3 mmol g−1 at 298 K and 0.2 bar,161 but MgO demonstrated only a 0.64 mmol g−1 uptake at 273 K and 0.2 bar.471 Hence, MgO generally has a limited potential for CO2 adsorption, especially in biogas separation. TiO2 (or titania), on the other hand, is a family of transition metal oxides consisting of several common polymorphs (i.e., rutile, anatase, and brookite) with tetragonal and orthorhombic crystal structures.472 Unlike MgO, the key advantages of TiO2 lie in its stability in water and the relative ease of obtaining well-dispersed nanoscale particles in solutions.473,474 Never-

3. COMPOSITE MEMBRANES 3.1. Definition and Motivations

Composite membranes (or MMMs) are defined as the integration of a dispersed filler (solid phase) into a continuous polymeric matrix.135 The concept behind composite membrane preparation is to leverage the key attributes of the fillers to alter the overall properties of the matrix materials and to obtain composite structures with enhanced gas transport properties. Primarily, the goal is to realize membrane performance that transcends the upper bound on the Robeson plot.135 This is typically achieved by exploiting either the selective nature of the filler to increase the membrane selectivity or the relatively higher diffusion coefficient of the gas solute in the filler (than in the polymer matrix) to obtain a higher gas permeability of the membrane. Previously, in section 2, we introduced some of the most important fillers for composite membranes to date and emphasized their inherent attributes, which are critical for promoting CO2 transport and improving CO2/CH4 separation. Herein, we focus on the various gas transport models and challenges encountered in integrating these fillers into the composite membranes. A glimpse of these challenges has already been provided in Table 17. Poor adhesion between the solid and polymer phase, which leads to nonideal interfacial morphology, is identified as a key challenge of composite membranes.476 To offer a better understanding of these challenges, we first consider mainstream mathematical models that are used for predicting the permeability of the more permeable gas component and describing the different nonidealities generated by the filler in the polymer matrix. Next, we tackle the root of the problem by discussing the fabrication techniques of composite membranes and share with readers some good practices for producing composite membranes with minimal defects. Lastly, a comparison and evaluation of the different type of interfacial defects are provided. Critically, the influence of the choice of polymer matrix on the formation of defects, the detailed evolution of specific types of interfacial defects during membrane formation, and the methodologies commonly adopted to overcome these defects are discussed. AG


crystalline microporous networks with metal or metal clusters linked to organic ligands via coordination bonding periodic arrangement of light elements via covalent bonding

MOFs

bulk carbonaceous materials with no particular morphological structures

materials with pore sizes between 2 to 50 nm

2D carbon nanosheets containing sp2 hybridized carbon arranged in honeycomb lattices

2D hydrated layered aluminosilicate minerals with interconnected tetrahedral and octahedral layers

carbon-based particles

mesoporous materials

GFMs

layered silicates

MOPs

crystalline aluminosilicates with interconnected SiO4 and AlO4 tetrahedra

compositional definition

zeolites

type

AH

external surface area: < 100 m2 g−1; pore size: 3−4 Å; pore volume: 0.03−0.4 cm3 g−1; high aspect ratio; hydrophilic; and crystalline

internal surface area: < 2000 m2 g−1; pore size: 20−500 Å; pore volume: 0.04−0.90 cm3 g−1; hydrophilic or hydrophobic; and crystalline or amorphous internal surface area: < 2000 m2 g−1; pore volume: 0.04−1.8 cm3 g−1; high aspect ratio; hydrophilic or hydrophobic; and crystalline

internal surface area: > 1000 m2 g−1; pore size: 5−30 Å; pore volume: 0.5 to >2 cm3 g−1; hydrophobic; crystalline or amorphous internal surface area: > 1000 m2 g−1; pore size: 2−20 Å; pore volume: 0.06−0.8 cm3 g−1; hydrophobic; and amorphous

internal surface area: < 1000 m2 g−1; pore size: 3−10 Å; hydrophilic (low Si/Al ratio) or hydrophobic (Si/Al = ∞); crystalline internal surface area: > 1000 m2 g−1; pore size: 4−15 Å; pore volume: 0.5 to >2 cm3 g−1; hydrophilic or hydrophobic; and crystalline

physicochemical property

single-layer graphene has little practical relevance; conflicting adsorption results owing to inconsistency in synthesis and functionalization; and functionalization is usually at the expense of structural integrity

single-layer graphene has lowest resistance given one atom thickness; strong chemical and thermal stability; impermeability creates tortuous path and suppresses CH4 permeability but improves CO2/CH4 selectivity; and readily functionalizable individual layers obtainable by exfoliation; high cationic exchange capacity; and swollen silicates possess appealing interlayer spacing for CO2/CH4 separation

hydrophilic surface reduce compatibility between filler and polymer and uniform exfoliation and orientation of layered silicate can be challenging

decreased crystallinity; large pores detrimental to overall selectivity; poor wetting results in poor compatibility with polymer matrix; and bonding between filler and polymer facilitates transport of gas molecules, leading to poor selectivity

tunable mesopore sizes by varying the alkyl chain length of the surfactant; enhanced permeability of composite membrane; and penetration of polymer matrix through the mesopore improves polymer−filler interaction

compatibility enhancement: silane coupling agents to enhance polymer−filler interaction and cationic-exchange (with surfactant) to allow effective tuning of the interlayer spacing; CO2 adsorption enhancement: silane coupling agents with amine groups

CO2 adsorption enhancement: surface functionalization of graphene surface with functional groups such as −NH2, −COOH, −OH, and −CH3, templates such as zeolites, MOFs, and MgO to increase surface area (>2000 m2 g−1), as well as enhance CO2 adsorption, and doping such as N- or S-doped GFMs

ACs and CMSs

rigid structures leading to nonselective voids within polymer matrix; moderate interaction with polymer matrix; and large particle size undesirable for use as filler

compatibility enhancement: “priming” on the filler surface with sizing agents (polymer wrapping), surface treatment with strong inorganic acids to give −OH and −COOH to promote dispersion and increase the overall free volume of polymer, surfactants to promote better dispersion in polymer matrix, and plasticizers (e.g., PEG) to assist the emulsification between filler (ACs) and the polymer; CO2 adsorption enhancement: postsynthetic functionalization such as amine grafting and metal cations impregnated onto CMSs to create a polarization effect compatibility enhancement: silane coupling agents to enhance polymer−filler interaction; CO2 adsorption enhancement: postsynthetic functionalization such as amine grafting and silane coupling agents with amine groups

hydrophobic with more resistance toward high humidity; strong chemical and thermal stability especially toward strong acids and bases; CMSs have narrower pore size distributions relative to ACs; and cost-effective

COF-10, CMP-5, PP-N-25, SNW-1, and PAF-1

CO2 adsorption enhancement: presynthetic functionalization with functional groups such as −OH, −NO2, −COOH, and −CH3 and postsynthetic functionalization such as −SO3H, amine

harsh synthesis conditions; poor scalability relative to zeolites and MOFs; lower CO2 uptake relative to zeolites and MOFs; and morphology control can be challenging

1:1 silicates (kaolin, serpentine) and 2:1 silicates (chlorites, smectites)

pristine graphene, GO, and rGO

MCM-41, MCM-48, SBA-15, and COK-12

Mg-MOF-74, HKUST-1, MOF-5, MIL-53, and ZIF-8

stability enhancement: azolate-based linkers (ZIF-series, formation of stronger metal−ligand bonding), hydrophobic functional groups; CO2 adsorption enhancement: presynthetic functionalization with functional groups such as −OH, −NO2, −CN, and −Cl, postsynthetic functionalization such as amine grafting, and coordinatively unsaturated open metal sites which possess high charge density (polarization effect to attract CO2)

susceptible to water hydrolysis; poor scalability relative to zeolites; and poor chemical and thermal stability

Ca-A (5A), DD3R, KMCM-22, Na-BEA, and SAPO-34

notable examples

greater versatility in modifying active sites for enhanced CO2 capture; tunable pore size with ligand modification; stronger polymer− filler compatibility relative to zeolites (due to presence of organic ligands); well-defined pores; and ease of morphological control strong polymer−filler compatibility; tunable pore size with ligand or monomer strut’s length modification; and highly functionalizable to give enhanced CO2 adsorption

compatibility enhancement: silane coupling agents to enhance polymer−filler interactions, low molecular weight materials (LMWMs), “priming” on filler’s surface, inorganic “whiskers” attached on the filler’s surface to increase surface roughness, and CO2 adsorption enhancement: postsynthetic functionalization such as amine grafting and ion-exchange with more electropositive cations

core strategy for improvement

strong water adsorption limits CO2 uptake; limited modification options for CO2 enhancement; and poor polymer−filler compatibility

main limitation

possess sites that allows selective uptake of CO2; well-defined pores; high scalability; strong chemical and thermal stability; and high cationic exchange capability

prominent merit

Table 17. Brief Overview of Emerging Materials with the Potential to Be Utilized As Effective Fillers in Composite Membranes for Biogas Upgrading



Chemical Reviews fumed silica, MgO, and TiO2 compatibility enhancement: silane functional agents to enhance polymer-filler interaction (fume silica, and TiO2); CO2 adsorption enhancement: silane functional groups such as (3-aminopropyl)trimethyoxysilane to improve affinity toward CO2 susceptible to water adsorption; high agglomeration tendency (non-linear expansion of overall free volume); and poor compatibility with polymer matrix improve membrane permeability by disrupting the polymer packing and increasing the polymer free volume

SWCNTs, DWCNTs, and MWCNTs compatibility enhancement: surface treatment via oxidation to give −OH and −COOH or functionalization with −NH2 groups to promote CNT dispersion in polymer matrix; CO2 adsorption enhancement: postsynthetic functionalization with silane or amine groups postsynthetic CNTs alignment remains challenging; agglomeration of CNTs in polymer matrix; and functionalization is usually at the expense of structural integrity

3.2. Mathematical Models of Gas Transport

One of the most important indicators to evaluate the performance of a gas separation membrane is the permeability, P. The permeability of a composite membrane is highly dependent on the transport and chemical properties of the polymer matrix, filler, as well as the intrinsic properties of the permeating gas species (i.e., size, shape, and polarity). Most gas separation membranes, whether symmetric or asymmetric, contain a dense layer to offer selective separation through a solution-diffusion mechanism (see section 1.3.1). The intricate interplay between the polymer matrix and filler is therefore critical in determining the diffusion rate of the permeating species through the membrane. This can be expressed by a diffusion coefficient, D, of the species. Critically, the transport properties of the polymer matrix and filler must be compatible to prevent one phase from exerting an overpowering mass transfer resistance, which can obscure the effect of the other phase, and dominate the overall permeability of the membrane.477 In addition, the chemical properties of the polymer matrix and filler dictate the interactions with the permeating species. This can be expressed via a solubility coefficient, S, of the species. Overall, the permeability of a particular gas species through a membrane can be conceptually described by the solution-diffusion mechanism as52

P=S×D

1D cylindrical carbon nanomaterials with nanoscale diameters

materials without intrinsic pores for gas transport

CNTs

nonporous materials

(1)

where P is permeability with units of barrer (1 barrer = 1 × 10−10 cm3 (STP) cm cm−2 s−1 cm Hg−1) while S and D are the solubility and diffusion coefficients, respectively. The permeability of a membrane can be experimentally determined using144 Q P = l AΔP

(2)

where l is the effective thickness of a membrane, Q is the volumetric flow rate of a particular gas through the membrane, A is the surface area of a membrane, and ΔP is the pressure drop across the membrane. To better grasp and quantify the effectiveness of a filler in the composite membrane, mathematical modeling is normally sought to give a more complete understanding of the change in transport properties after the filler is incorporated. Modeling is also an extremely useful tool in determining the necessary filler morphological structures and the critical filler loadings to obtain optimized gas separation performance of the composite membranes.478,479 The Maxwell model is one of the most classical models, which was originally developed based on the electrical transport of composite materials in the presence of dielectric medium as shown below:

internal surface area: < 2000 m2 g−1; pore size: 4−150 Å; pore volume: 0.3−0.8 cm3 g−1; high aspect ratio; hydrophilic or hydrophobic; and crystalline external surface area: 20 vol %; and ignores the effect of filler shape, size distribution, and agglomeration

Bruggeman

electrical transport

Maxwellb

requires trial-and-error; calculation of gas transport properties at ϕm is not feasible; and ignores the effect of filler shape, size distribution and agglomeration

deviates less as compared to Maxwell

spherical generally larger than 20 vol % but smaller than ϕm

implicit

dielectric constant

Böttcher

requires trial-and-error; calculation of gas transport properties at ϕm is not feasible; and ignores the effect of filler shape, size distribution, and agglomeration

deviates less as compared to Maxwell

spherical generally larger than 20 vol % but smaller than ϕm

implicit

dielectric constant

Higuchi

can account for the effect of particle shape, size distribution, and agglomerationg and gas transport properties can be calculated explicitly displays diverging behavior when Pf/Pp → ∞

any shape 0 < ϕf < ϕm d

explicit

elastic modulus

Lewis-Nielsen

can account for the effect of particle shape, size distribution, and agglomerationg requires trial-anderror

any shape 0 < ϕf < ϕm e

implicit

thermal conductivity

Pal

restricts to mathematical prediction only

gas transport properties can be calculated explicitly

lamellar −f

explicit

−

Cusslerc

a The modeling equation summarized to date does not include the modeling of nonporous nanoparticle. bMaxwell equation is part of the Maxwell−Wagner-Sillars equation at n = 1/3.481 cThis model is not an adaptation of any existing model. It is modeled based on the diffusion of gas molecules in flake-containing membranes. dAs ϕm → 1, the expression is reduced to the Maxwell equation. eAs ϕm → 1, the expression is reduced to the Bruggeman equation. fNo information on the range of filler loading. gThe model predicts accurate behavior when ϕf → ϕm. This allows structural parameters to be considered in the model with a known value of ϕm for a specific filler.494

main limitation

physical basis (origin) function form filler shape filler loading range, ϕf (vol %) key merit

parameter

Table 18. Comparison of Representative Gas Transport Models for Ideal Composite Membranes485−487,491−493a


AJ


AK

works only for low filler loading ratios; calculation of gas transport properties at ϕm is not feasible; ignores the effect of particle shape, size distribution, and agglomeration; and fails when considering simultaneously two or more nonidealities

works only for low filler loading ratios; calculation of gas transport properties at ϕm is not feasible; and ignores the effect of particle shape, size distribution, and agglomeration

sieve-in-a-cage; plugged sieves; and rigidified interface

all nonideal morphologies can be considered each time

works only for low filler loading ratios; does not allow simultaneous consideration of two or more nonidealities; calculation of gas transport properties at ϕm is not feasible; and ignores the effect of particle shape, size distribution, and agglomeration

sieve-in-a-cage

the effect of particle shape on overall permeability can be expressed

works only for low filler loading ratios; requires trial-anderror; and ignores the effect of particle size distribution and agglomeration

simpler computation than modifiedMaxwell (Mahajan)

sieve-in-a-cage; plugged sieves; and rigidified interface more than one nonidealities can be simultaneously considered in the model when calculating the overall permeability

spherical 0 < ϕf < 20

spherical 0 < ϕf < 20

sieve-in-a-cage; plugged sieves; and rigidified interface

spherical 0 < ϕf < 20e

explicit

any shaped 0 < ϕf < 20

explicit

explicit


Felske

implicit

electric transport

Li

electric transport

Mahajan (modified-Maxwell)

effective medium theory

Erdem-Ş enatalar

requires extensive precomputation before the effective permeability can be determined and fails when considering simultaneously two or more nonidealities

sieve-in-a-cage; plugged sieves; and rigidified interface possible to consider the effect of particle shape, size distribution, and agglomerationh

spherical 0 < ϕf < ϕmf

explicit


Pal (modified-Felske)

Hashemifard

applicable for ideal composite membrane when interfacial effects are ignored

possible to consider the effect of particle shape, size distribution, and agglomerationh cannot model sieve-in-a-cage and plugged sieve nonidealities and requires trialand-error

requires extensive precomputation before determining the effective permeability; cannot model plugged sieve nonideality; and ignores the effect of particle shape, size distribution, and agglomeration

sieve-in-a-cage and rigidified interface

spherical high particle loadingsg

explicit

−c

rigidified interface

any shaped 0 < ϕf < ϕm

implicit


Shimekit (modified-Pal)b

a The modeling equations summarized to date do not include the equations for nonporous nanoparticles. bThe modification is based on the calculated permeability of ideal composite membrane provided by Pal.499 cThis model is not an adaptation of any existing model. It is described based on the flow pattern of gas molecules in composite membranes. dIt requires the input of characteristic length for better prediction of the effective permeability. eThe equation is reduced to the Maxwell equation if interfacial layer is absent in composite membrane. fAs ϕm → 1, the expression is reduced to the Felske equation. If interfacial layer is absent, the expression is reduced to the Lewis-Nielson equation. Considering the absence of interfacial layer and ϕm → 1, the expression is reduced to the Maxwell equation. gNo information on whether the model holds at maximum volume fraction of fillers, ϕm. hStructural parameters can be considered in the model with a known value of ϕm for a specific filler.

main limitation

physical basis (origin) function form filler shape filler loading range, ϕf (vol %) type of nonideal interface key merit

parameter

Table 19. Comparison of Representative Gas Transport Models for Non-Ideal Composite Membranes479,495−501a



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In the Maxwell model, the gas transport profile around a filler is not affected by the presence of other fillers. On the basis of this assumption, the model is strictly relevant only at low filler loading of volume fraction less than 0.2.480,481 Above this value, an onset of deviation (from eq 3) exists, which can be defined as a percolation threshold.57 The mathematical concept of the percolation threshold, though not discussed here, is covered elsewhere (see refs 480−483). Beyond this threshold limit, highly interconnected channels between fillers are formed due to agglomeration and this leads to a complete breakdown of the model, resulting in large deviations from the experimental results. The maximum packing volume fraction of fillers, ϕm, is often stipulated as 0.64 based on random closed packing of uniform spheres. Also, the Maxwell model assumes a homogeneous dispersion of the fillers with seamless interfacial contact between the filler and polymer matrix. Fillers for composite membranes hitherto exhibit a wide variety of dimensional and morphological structures, along with different sizes and distributions (see section 2). In this regard, the Maxwell model is oversimplified and fails to consider the complex structural parameters of the fillers.480,481 Several modifications have since been made to augment the model. The Bruggeman model incorporates a wider range of ϕf than that defined by the Maxwell model.484 The Lewis-Nielsen model, on the other hand, takes into account the effect of filler morphology on the permeability of composite membranes through a factor, ψ:485 Pcm = Pp

where G is a geometric factor that accounts for the dimensional structure of the fillers. For 1D fillers (such as CNTs) that are transverse to the gas flow direction, G is equal to 1. For spherical or isometric particles, G is 2, and for lamellar fillers, G takes a value between 0 to infinity, depending on the orientation of the lamellar fillers with respect to the gas flow direction. The aforementioned models are applicable only to composite membranes with ideal interfacial morphologies (Table 18). They are often inadequate in predicting the gas transport properties of membranes with nonideal interfacial morphologies, which are in fact more representative of composite membranes in reality.489 This inadequacy largely arises from the assumption of linear concentration profile across the membrane and the embedment of the filler properties into a single empirical morphology-related parameter. By doing so, the permeability of the composite membranes becomes independent of the membrane thickness as well as interphase thickness that is driven by the filler particle size. When this happens, the models are incapable of reconciling with experimental results. Hence, nonidealities need to be postulated to correct for the inconsistencies.490 This can be achieved through an extensive process of modifying the models, which involves multiple rounds of derivation, trial-and-error, and validation with empirical experimental data, to include specific interfacial regions at the polymer−filler interface. A perfected model is often identified by a strong agreement between the predicted and experimental data, and this provides solid evidence of a specific interfacial morphology in the composite membrane. Despite this, few modified transport models are versatile enough to simultaneously consider multiple nonideal interfacial morphologies alongside fillers of different phases, shapes, sizes, and size distributions (Table 19).481 Nevertheless, membrane permeation and filler sorption experimental data, coupled with the most congruent modified transport model, is still (at present) regarded as the most powerful method to probe the polymer− filler interfacial morphology of less complex composite membranes. This method is also incentivized by the challenge in characterizing the volume fraction of the interfaces using conventional microscopic techniques.479 In view of this, we have assembled and summarized representative models and the critical parameters used to describe ideal and nonideal composite membranes (Table 18 and Table 19). These tables are constructed in a concise manner to showcase models that are pertinent for nanoporous materials. For a more comprehensive discussion on this subject, interested readers can refer to a review by Hoang et al. (see ref 481).

1 + 2ϕf (α − 1)/(α + 2) 1 − ψϕf (α − 1)/(α + 2)

ij 1 − ϕ yz mz zzϕ where ψ = 1 + jjjj j ϕ 2 zz f k m {

(4)

(5)

The ϕm is correlated to the particle shape, size distribution, and the aggregation state of the fillers. As such, since ψ is a function of ϕm, the Lewis-Nielsen model takes the filler morphology into consideration when predicting the permeability of composite membranes. It is also noteworthy to mention that as ϕm → 1, eq 4 simplifies to the Maxwell equation (eq 3). The Cussler model is more specific and accounts for composite membranes with lamellar fillers. Its equation has a similar form to that of the Maxwell model:486 Pcm = Pp

1 (1 − ϕ ) y ij Pf 1 − ϕf + 1/jj ϕ P + 4 2 2f zzz αf ϕf p f k {

(6)

3.3. Membranes and Modules

where αf is the filler aspect ratio and ϕf is the volume fraction of the lamellar fillers. 2D fillers such as 2D MOFs, 2D COFs, GFMs, and layered silicates are best described by αf and thus the Cussler model is well-suited for composite membranes with 2D fillers. Beyond 2D fillers, a model proposed by Petropoulous487 and later extended by Toy and co-workers488 allow fillers of different dimensional and morphological structures to be taken into consideration: ij yz jj zz jj zz (1 G ) + ϕ f j zz Pcm = Ppjj1 + zz jj P / P + G ij f p zy zz jj − ϕ j z j Pf / Pp − 1 z j z fz k { k {

3.3.1. Membrane Fabrication Techniques. Polymeric gas separation membranes are mostly synthesized via a phase inversion method. The simplest technique for preparing flat sheet phase inversion membranes is by solution casting. Dense symmetrical membranes are typically made by first dissolving a polymer in a solvent system (single- or mixed-solvent).40 Then, the polymer dope solution is cast onto a smooth support (e.g., a glass plate) via a doctor blade technique before enclosing it in an atmosphere (usually inert) to allow controlled solvent evaporation.502 In academic research, fabrication of composite membranes follows a largely similar technique except that the fillers are directly blended into the dope solution before membrane casting.480 Despite its facile

(7) AL


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1970’s was the breakthrough in producing low-cost membrane modules.57 For this reason, we cannot emphasize enough the importance of membrane modules on the path toward successful commercialization of composite membranes. Unquestionably, the development of membrane modules for composite membranes must grow in tandem with the membrane fabrication techniques. However, research efforts in this area are not always as visible as compared to efforts spent on synthesizing high-performance composite membranes. This is because module technologies and their knowhow are often trade secrets in the membrane industry, with basic details revealed only in patents. Hence, many academic researchers have little access to such technical knowledge. Three main types of membrane modules are used in gas separation: (1) plate-and-frame, (2) spiral wound, and (3) hollow fiber modules.64 The plate-and-frame modules are one of the earliest membrane systems, which utilize simple platelike frames to support the flat sheet membranes. Today, plate-andframe modules have been superseded mostly by spiral wound and hollow fiber modules due to more competitive production costs and smaller footprints (Table 20). Apart from that, the

implementation, good practices must still be observed to ensure successful fabrications of composite membranes. First, aggregated fillers must be homogeneously dispersed in a suitable solvent before mixing into the polymer dope solution. This is realized either by bath-sonication or applying a higher-intensity probe-sonication for more tenacious aggregation. Centrifugation after sonication and using the filler left in the supernatant solution can further reduce the presence of large aggregates. Importantly, dispersed fillers suspended in solvents are in a continuous state of reaggregation when left to stand. We recommend standardizing the standing duration of the dispersed filler solution to maintain a similar aggregation state of the filler and achieve consistent structural integrity of the composite membrane. Second, the polymer dope solution must possess a suitable viscosity to ensure smooth and uniform casting. A polymer concentration of ∼15−20 wt % usually suffices to give a casting solution with an acceptable viscosity.57 For composite membrane fabrication, increasing the viscosity is useful in reducing sedimentation of fillers at the solvent evaporation step.476,503 Third, controlling the rate of solvent evaporation is critical. Poor control of the evaporation rate can result in a nonideal polymer−filler interface, especially for fillers such as TiO2 and SiO2 that are generated in situ within the polymer matrix during phase inversion.504 Fourth, annealing the ensuing membrane at a temperature above the Tg of the polymer matrix can help to heal defects that arise during membrane fabrication. The process of annealing not only enhances the flexibility of the polymer chains but also establishes a more intimate contact with the filler and facilitates a stronger polymer−filler interaction.505 However, the use of high annealing temperatures is not always suitable, considering the thermal stability of specific fillers. Additional methodologies to improve interfacial interactions will be described in section 3.4.2. Hollow fiber membranes, on the other hand, are prepared via one of four main approaches which include melt-, dry-, dryjet wet-, and wet-spinning method.502 In particular, for dry-jet wet- and wet-spinning, phase inversion is driven by immersion precipitation, where the polymer dope is extruded into a coagulation bath containing a nonsolvent. Precipitation, in this case, arises due to an exchange between the solvent and nonsolvent. By and large, the above good practices stay relevant for hollow fiber composite membranes except for the solvent evaporation step, which does not apply to the wetspinning of hollow fibers. For TFC (or multilayered) gas separation membranes, the dense selective layer is usually prepared by dip- or spin-coating, interfacial polymerization (IP) or solution-casting.506 As the layer is designed to be thin relative to the substrate (typically between 0.1−1.0 μm), integrating fillers into this selective layer presents a greater challenge.55 Limitations on the particle size and distribution, as well as the state of dispersion of the filler, become increasingly significant with decreasing selective layer thickness. It is postulated that a 100 nm thick selective layer will necessitate reducing the size of the filler particle to the order of 10 nm to preserve the integrity of the layer.135 A filler size of this magnitude implies that other issues such as morphological structure, dispersion stability, and particle size distribution of the filler could become pivotal in a defect-free selective layer.135 3.3.2. Module Designs. One of the reasons behind the commercialization of membrane processes in the 1960’s and

Table 20. Overview of the Three Main Module Designs for Gas Separation Membranes57,63,64 specification

hollow fiber

spiral wound

plate-and-frame

production cost (US$ m−2) packing density (m2 m−3) approximate membrane area per module (m2) pressure drop (permeate side) capacity for high-pressure operation capacity for concentration polarization control capacity to encompass all membrane materials potential for composite membranes

5−20 500−10000 300−600

5−100 200−1000 20−40

50−200 30−500 5−20

high

moderate

low

high

high

high

low

moderate

high

limited

high

high

limiteda

high

high

a

Limited potential is based on current technical challenges in spinning and modifying composite hollow fiber membranes and the limited number of polymeric materials which can be spun into hollow fibers. Furthermore, the low production cost is driven by high-volume manufacturing to justify the massive investment in the module assembly equipment. Fabricating hollow fiber composite membranes at this scale is currently unachievable.

spiral wound and hollow fiber modules also exhibit other advantages. A brief overview of these advantages is summarized in Table 20. For a more comprehensive discussion, readers can refer to earlier reviews (see refs 57 and 64). At present, there are no modules specially designed for composite membranes due to their lack of commercial acceptance. As such, module designs for composite membranes are likely to replicate that of the conventional polymeric membranes. For academic research, considering the ease of fabricating and modifying flat sheet composite membranes (see section 3.3.1), flat sheet modules are more commonly utilized. Hence, we anticipate that spiral wound modules may have a lower market entry barrier than hollow fiber modules (Table 20). Nevertheless, future efforts should still focus on assembly methods, optimizing gas flow distribution and pressure drop within the modules as well as securing flawless sealing and AM


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Figure 29. (a) Impact of ideal (following prediction from Maxwell Equation) and nonideal morphologies on the performances of composite membranes. Adapted with permission from ref 479. Copyright 2005 Elsevier. (b) Signature CO2 transport profiles of various interfacial morphologies of composite membranes. The normal profile refers to the diffusivity of CO2 molecules in the polymer phase.

in the free volume. The polymer chains, at this point, remain flexible and can undergo relaxations to reach their new equilibrium conformations.508 However, as more solvent continues to evaporate, the Tg reaches the solvent evaporation temperature, and vitrification is initiated at this point. The polymer then becomes glassy, and the relaxation time via polymer diffusion becomes extremely long. When this happens, it is impossible for the polymer chains to diffuse freely into one another and respond adaptively to a potential stress. On the contrary, the filler is a solid phase that is less affected by solvent evaporation unlike the polymer matrix. This is because most fillers are relatively more rigid and experience little or no change in the volume and shape. Hence, as solvent evaporates, there exists an uneven distribution of stress, especially at the polymer−filler interface.479 Correspondingly, those chain of events leads to the formation of the sieve-in-a-cage and rigidified nonideal interfacial morphologies. 3.4.1.1. Sieve-in-a-Cage. The sieve-in-a-cage nonideal morphology is a direct consequence of the uneven distribution of stress in composite membranes. Essentially, the stress stems from the inability of the polymer chains to relax in all directions because of the infinitesimal volume reduction of the rigid fillers during the evaporation of solvent. This is exacerbated by the fact that the polymer−filler interaction is typically weaker than the polymer−polymer interaction. In particular, if the integrated filler is inorganic (hydrophilic) in nature, it brings about an organic−inorganic incompatibility

potting of membranes, considering a change in the physicochemical properties of the surface of composite membranes.64 Particularly, modeling of different module configurations, flow patterns, and operating conditions will help advance module designs for composite membranes with higher permeabilities and/or selectivities after the incorporation of filler materials.64,507 3.4. Nonideal Interfacial Morphologies

The greatest repercussion of nonideal interfacial morphologies is the loss in separation performances of the composite membranes. Therefore, in this section, we aim to address this problem by first understanding the origin of each type of nonideal morphologies and assessing their impacts on the CO2/CH4 membrane performances. Next, strategies to mitigate these nonidealities are taken into account. For composite gas separation membranes, the nonideal morphologies of interest are sieve-in-a-cage, rigidified interface, and plugged sieves. As a prelude, these morphologies can be perceived as shown in Figure 29. In addition, their impacts and signature CO2 transport profiles are illustrated and compared against the ideal morphology. 3.4.1. Origins of Defects. The formation of nonideal interfacial morphologies is closely related to the development of stress during the vitrification process. As the solvent evaporates from a cast film, the viscosity of the dope solution increases, which results in an increase in the Tg and a decrease AN


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Figure 30. Comparison between (a) sieve-in-a-cage and (b) rigidified interface formation in composite membrane. Reprinted with permission from ref 479. Copyright 2005 Elsevier.

mitigate the sieve-in-a-cage morphology. This will be elaborated further in section 3.4.2.1. 3.4.1.2. Rigidified Interface. Once the vitrification process initiates, the diffusion of polymer chains over the time scale of solvent evaporation becomes almost negligible. In the case of a purely polymeric film, the polymer matrix continues to contract isotropically with solvent evaporation beyond the vitrification point. For a composite membrane, the contraction occurs toward the filler surface and gets inhibited owing to the rigidity of the filler.509 The result is a compressive stress at the polymer−filler interface where polymer chains pile up at the filler surface and extend many layers into the bulk polymer, leading to a condensed interface that is rigidified (Figure 30b).479 As a corroboration of this hypothesis, Moore and coworkers integrated 5 μm zeolite 4A and 1 μm CMSs into Matrimid membranes and found that the rigidified thickness was ∼0.6−0.9 and 0.05 μm, respectively. They concluded that the effect of inhibited contraction was less with a smaller filler size, providing partial evidence of inhibited contraction as the cause of rigidified interface.509 Apart from inhibited contraction, the use of silane coupling agents on the surface of fillers has also been associated with the formation of the rigidified interface. However, unlike the sieve-in-a-cage morphology, the silane groups in this case exert an exceedingly strong interaction with the polymer such that the polymer chain mobility is irrevocably inhibited to induce a compressive stress and render the polymer−filler interface rigidified.510 The gas transport profile in the rigidified interface is distinctively different as compared to that of the sieve-in-a-cage morphology (Figure 29b). Specifically, in the rigidified interface, solution-diffusion transport continues to remain as the dominant mechanism because of the reduced fractional free volume (FFV) of the condensed polymer interface. This promotes a decrease in the diffusivity of the gas molecules and thus lowers the gas permeability. In fact, the permeability in this rigidified interface can be up to three to four times smaller than that of the reference membranes as modeled by a study by Moore and Koros.479 Such a reduction has a profound implication on the separation performance of the composite membranes. This is particularly so for membranes with high filler loadings where the number of rigidified regions grows extensively within the polymer matrix. In this instance, the decline in gas permeability in the rigidified interface can potentially offset the higher gas diffusion through the pores of the filler, leading to an overall reduction in the membrane

issue, which, under the substantial stress generated during relaxing, can emerge as a failure at the polymer−filler interface.135 For that reason, the polymer matrix adheres poorly to the filler surface, giving rise to the delamination and formation of interfacial voids around the fillers (Figure 30a). Agglomeration is another common occurrence that results in the formation of the sieve-in-a-cage morphology. When the filler agglomerates, the shape deviates while the size increases in a random manner to incur a wide particle size distribution. This weakens the interaction with the polymer matrix and increases the likelihood of defective interfacial voids around the agglomerates. Additionally, agglomeration of the filler is also responsible for the formation of defective pinholes in the skin layers of the ISA and TFC membranes. The mitigation of such sieve-in-a-cage nonideal morphology is particularly challenging. Even annealing at a temperature above the Tg of the polymer does not always guarantee success (see section 3.3), owing to the huge difference in the thermal expansion coefficients between the polymer matrix and filler.479 A high free volume region is induced by the presence of interfacial void between the polymer matrix and filler. Generally, the dimension of this region is larger than the pore size of the filler. Hence, the interfacial void serves as a supplementary gas transport channel, which causes a deviation in the transport profile of the gas molecules. By and large, the gas molecules diffuse via a solution-diffusion mechanism through the dense polymer matrix. However, in the interfacial void, the transport mechanism is often predominated by Knudsen-like diffusion, which has lower mass transfer resistance but poorer gas separation selectivity. Therefore, gas molecules are most likely to diffuse along the interfacial void rather than through the polymer matrix or pores of the filler (Figure 29b). As a result, composite membranes with sieve-in-a-cage morphology tend to exhibit a trade-off in their separation performances, which see an increase in gas permeability at the expense of membrane selectivity. In the case where the effective void thickness or high free volume region is similar to the size of the gas molecule, the sieve-in-acage morphology also results in an enhancement in the gas permeability without compromising the membrane selectivity.481 However, precise engineering of the void thickness is extremely difficult to achieve. Therefore, other simpler approaches such as the use of silane coupling agents to improve the polymer−filler adhesion have been adopted to AO


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Figure 31. (a) The effect of Mg2(dobdc) in different polymer matrices and the CO2 uptake profiles of Mg2(dobdc) in rubbery polymers (b) PDMS, (c) XLPEO, as well as glassy polymer, and (d) polyimide (PI). The solid green and blue symbols indicate the profile of the unfilled and composite membranes, respectively, while the open symbol shows the theoretical maximum. Adapted with permission from ref 513. Copyright 2013 Royal Society of Chemistry.

permeability (Figure 29b). Nevertheless, the membrane selectivity is usually enhanced at the same time due to the condensed polymer interface which offers greater selectivity.511 Thus, it is still possible for membranes with rigidified interface to deliver competitive separation performances that surpass the Robeson upper bound limit, especially when polymers of high intrinsic permeability are exploited as the matrix materials (Figure 29a). 3.4.1.3. Plugged Sieves. In contrast to the sieve-in-a-cage morphology and rigidified interface, plugged sieves emerge because of pore blockage rather than filler-induced stress on the matrix (polymer chains). The pores of fillers are susceptible to plugging by water, solvents, contaminants, minor components that can be present in the feed gas,479 and even more so by the flexible polymer chains during membrane fabrications.503,511 Pores can be totally or partially plugged depending on the extent of pore blockage. If the pores are totally plugged, the gas molecules are incapable of diffusing through the pores of the filler. In this case, the plugged fillers behave like nonporous fillers and the membrane selectivity is usually not enhanced, assuming that the fillers do not alter the chain packing and free volume of the polymer matrix.512 On the other hand, when the pores are partially plugged, the composite membrane generally shows a drop in the gas permeability as compared to when the pores are completely unplugged (Figure 29b). The drop is often accompanied by a change in the membrane selectivity. Essentially, the pore size of the filler and the dimension of the gas molecule are key drivers affecting the membrane selectivity. When the original pore size

is in the range of the diameter of the gas molecule, the membrane selectivity is reduced with pore plugging. Otherwise, the selectivity is likely to enhance when the pore size is much larger than the diameter of the gas molecule.481,511 In the latter case, pore plugging helps to narrow the pore size and its distribution, rendering the filler more selective and the composite membranes with greater capacities to deliver higher separation performances. The extent of pore plugging is also closely related to the type of polymer matrix used (Figure 31a). As demonstrated by Bae and Long,513 the plugged sieve morphology was more prominent when rubbery polymers, such as cross-linked PEO (XLPEO) and PDMS, were used to encapsulate the Mg2(dobdc) filler. This was evident by the significantly lower CO2 uptake profiles of the composite membranes as compared to the theoretical predictions (Figure 31b,c). When a glassy polymer like polyimide was utilized instead, the CO2 uptake has a similar profile to the theoretical calculations, suggesting that the pores of the Mg2(dobdc) fillers remained effectively available after encapsulation (Figure 31d). The authors attributed this difference to the molecular structure and high mobility of the rubbery polymer chains. Also, the pore size of the filler is a critical factor, given that pore plugging by rubbery polymer chains was found for Mg2(dobdc) with a pore size greater than 1 nm but not for fillers, including zeolites, SIFSIX-3-Zn and UiO-66, which possessed pore sizes smaller than 1 nm.106,513,514 It is noteworthy to mention here that for low-pressure CO2 (ad)sorption measurements, chain rigidification can lead to slow diffusion and long equilibration AP


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Figure 32. (a) Examples of common silane coupling agents and (b) schemes showing the reaction mechanism of silane coupling agent at the polymer−filler interface. Adapted with permission from ref 517. Copyright 2008 Elsevier.

filler. A mismatch, where a smaller coupling agent is used together with a filler of larger pore size, can instead be counterproductive by resulting in undesirable plugging of the pores. In view of this, a more comprehensive understanding of each mitigating strategy is essential for its successful implementation. This will be provided in the upcoming sections. 3.4.2.1. Silane Coupling Agents. Using silane coupling agents to modify the surface of fillers is by far the most common approach to enhance the adhesion between the filler and polymer matrix. Currently, this method is used extensively for the modification of zeolites and silica surfaces, owing to the presence of hydroxyl groups on the fillers. In fact, the method is so versatile that it can be applied to other fillers as long as they contain surface hydroxyl groups to react with the coupling agents via facile silane chemistry. Typically, a silane coupling agent comprises two different reactive groups on its silicon atom (Figure 32a): (1) the alkoxide groups, which hydrolyze readily in water to give silanol groups that can interact with hydroxyl groups on the surface of the filler and (2) an organofunctional group such as amino, epoxy, or methacryloxy that is capable of extending and interacting with the polymer chains.516 Due to this unique structure, the silane coupling agents are able to tether together materials of a very different nature to help promote a strong polymer−filler interaction. The reaction mechanisms of modifying the filler surface and forming the polymer−filler interaction are described in Figure 32b. Briefly, the modification first involves a hydrolysis reaction to yield silanol groups on the coupling agent, which then condense with one another to give very stable siloxane Si−O−Si bonds. The remaining silanol groups can act as coupling points, undergoing the same condensation reaction

time. Researchers sometimes do not recognize this fact and fail to provide sufficient time for the isotherm to reach equilibrium. As a result, low CO2 uptakes are often misinterpreted as due to pore plugging. Hence, for the credibility of the (ad)sorption results, we encourage researchers to allow enough time to elapse so that equilibrium can be reached during measurements. 3.4.2. Mitigation Strategies. Due to the disparate origins of the nonidealities, different strategies are called upon to mitigate different types of interfacial morphologies. At present, the proposed strategies are largely designed to maintain the flexibility of polymer chains during and after the membrane preparation, as well as to promote interfacial interaction and adhesion of the filler with the polymer matrices.481 They are pursued through various means such as thermal annealing, membrane preparation via melt processing, addition of plasticizer, and low molecular weight materials (LMWMs) into the polymer dope solution, using glassy−rubbery copolymer as doping polymer and surface modification of the filler.480 In principle, these efforts are intended at releasing the interfacial stresses and creating defect-free interfaces. Therefore, they are well-suited for mitigating the sieve-in-acage and rigidified interfacial nonidealities. Even more so, they are effective for resolving the inherent organic−inorganic incompatibility issue as in the case with zeolite fillers.515 On the contrary, for plugged sieves morphology, the mitigation steps are usually put in place to remove water, contaminants, solvents, or other components which are capable of plugging the pores of the fillers. Alternatively, the pores can be protected by modifying the surface of the fillers with coupling agents that are overly large in molecular dimensions. However, this strategy requires careful matching of the coupling agent and AQ


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interfacial void between the filler and polymer matrix (Figure 33).522

with the surface hydroxyl groups of the filler to afford the silane-modified filler. Successively, the organo-functional group extends toward the polymer chains to chemically react and create the covalent bonds necessary for the polymer−filler interaction (Figure 32b).497,517 As a result of this ability to form covalent bonds on both ends, the silane coupling agents can act as effective compatibilizers to help adhere the interface between the solid (filler) and polymer phase, making them ideal for mitigating the sieve-in-a-cage interfacial morphology. Despite the advantages of silane-coupling agents, careful optimizations have to be made for successful implementation of this strategy because poor microscopic adhesions (as a result of imperfect closure) may exist persistently even though the polymer−filler interface appears macroscopically well-adhered after the silane modification.518 When this happens, the defective interface is not eliminated but simply reduces to a nanometric void with a higher mass transfer resistance than the sieve-in-a-cage morphology.510,519 This reduced void, however, stays large enough for gas molecules to leak through nonselectively. Aptly termed as a leaky interface by Mahajan and Koros, poorer transport properties of the composite membranes therefore follow, leading to decreased gas permeability without any improvement in the membrane selectivity. To make matters worse, a decline in selectivity is also commonly observed, especially if the polymer matrix used is rigid like in the case of Matrimid.479,519 In some other instances, the addition of silane coupling agents can counterintuitively induce defect formation. For example, coupling agents such as APTMS (Figure 32a) are known to plug the pores of fillers given the large amount of coupling points and small molecular dimensions, which facilitate easy entry into the pores.429 Thus, to avoid formation of such defects, APDEMS is preferred, owing to the fact that it is larger in dimension. Moreover, it consists of two ethoxy and one methyl groups on the silicon atom (Figure 32a). Compared to APTES with three ethoxy groups, APDEMS offers less coupling points upon hydrolysis and this reduces the likelihood of blocking the pores of the fillers.497 Again, the example shows that it is imperative to match the filler with the right silane coupling agent for optimization purposes. Critically, optimizing the threshold loading of silane coupling agent is also essential for minimizing interfacial defects, and achieving the desired separation performances of the composite membranes.518,520 3.4.2.2. Low Molecular Weight Materials. Analogous to the silane coupling agents, LMWMs are likewise able to create the tethering effect necessary for closing the polymer−filler interface. LMWMs are long aliphatic or polyaromatic compounds with polar functional groups to interact with the polymer matrix and modify its chain characteristics.521 The main advantage of LMWMs is that they can be introduced as additives into the polymer dope solutions together with the fillers. Hence, employing LMWMs for mitigating the sieve-ina-cage morphology is a much more convenient and easier to implement strategy, considering that the silane coupling agent needs to be first reacted onto the filler surface before incorporating the modified filler into the polymer matrix. In principle, LMWMs play the same role of a compatibilizer or as a third component through which the polymer−filler interaction is established. However, instead of forming covalent bonds, strong intermolecular forces of attraction are used to promote interfacial adhesion in this case. This is exemplified by LMWMs that utilize hydroxyl, amino, or nitro groups to interact via hydrogen bonding and close the

Figure 33. Scheme showing a LMWM interacting via hydrogen bonding at the interface of the filler and polymer matrix. Adapted with permission from ref 523. Copyright 2001 Elsevier.

Although the incorporation of LMWMs into polymer matrices as additives typically increases the selectivity of composite membranes, they often carry a crucial drawback, which is a reduction in the gas permeability.523 This response is consistent with an effect commonly known as antiplasticization.524 Through the interaction of LMWMs with the polymer chains, the overall stiffness in the polymer matrix can increase to reduce the segmental motion of the polymer chains, as well as the free volume of the matrix.116,521,525 Particularly, when LMWMs are added in high concentration, the antiplasticization effect can dominate to the extent that it deteriorates the gas transport properties of the polymer phase and brings about detrimental effects to the overall separation performance of the composite membranes.505,522 Thus, without careful optimization of the added amount of LMWMs, the success rate of this mitigating strategy is low. In addition, the choice of LMWMs is another limiting parameter. Ideally, LMWMs should possess sufficiently high molecular weights to impede loss through evaporation during membrane fabrications but at the same time have low enough molecular weights to enable them to remain soluble in the solvents used for making the polymer dope solutions. For this reason, finding a delicate balance in the molecular weight is crucial in ensuring high miscibility and low volatility of the additives. Coupled with the right polarity of the functional groups, the selected LMWMs can then fulfill their roles by intercalating within the interfacial void to promote strong interaction between the filler and polymer matrix. 3.4.2.3. Surface Modifications of Fillers. Apart from the silane coupling agents, there are some other ways of modifying the surface of fillers to promote the polymer−filler interaction and achieve enhanced interfacial adhesion. Physical deposition or coating of the filler surfaces with polymers is one such example. By and large, such a modification involves either coating of the dope polymer, or a nonidentical polymer, to create a thin layer of polymer on the surface of the filler. In the former case, the method is known as “priming” while in the latter, it is known as “sizing”. The general protocol to carry out this surface modification is summarized in Figure 34. Priming is a technique used to promote interfacial adhesion between the polymer and filler, as well as increase filler loadings in composite membranes.480,481 The strategy was first adopted by Mahajan as a protocol to increase the zeolite 4A AR


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Figure 34. Scheme showing a general protocol of surface modification of the filler prior to incorporation into the polymer matrix.476,496

loadings in the composite membranes, as the filler was prone to strong agglomeration at high loadings.476 Agglomeration of fillers makes the fabrication of composite membranes challenging given the possibility of defects such as the sievein-a-cage morphology. Priming, therefore, allows the dope polymer to wrap around the fillers in their dispersed state and helps retard their rate of agglomeration during the fabrication process.496 This technique is also effective for nanosized fillers, which are generally less stable and prone to more severe agglomeration than submicron fillers (∼0.3−5 μm).476 Hence, by alleviating agglomeration and having a compatible polymer layer on the filler surface to exert strong adhesion with the bulk polymer matrix, defect-free composite membranes with greater versatility to accommodate higher filler loading and smaller filler particle size can be realized. However, to ensure the success of this strategy, selecting a strong nonsolvent for priming of the filler is important because the polymer chains on the filler are unable to vitrify entirely when a weak nonsolvent is used to precipitate the primed-filler (Figure 34). Consequently, the polymer chains remain capable of entangling and bridging with polymer chains from another adjacent filler, leading to a phenomenon called “conglomeration” where multiple primed-fillers cluster together to reduce the effectiveness of the priming strategy.496 Sizing is another similar technique which employs polymers, known as sizing agents, to modify the filler surfaces as opposed to using the dope polymers (as in the case of priming). The sizing agent value-adds the filler by increasing the number of chemical reactive sites and/or surface area for interaction with the polymer matrix.480,481 In addition, the polymer chains of the sizing layer can get entangled with the chains of the polymer matrix at the interface to enhance the polymer−filler adhesion via an interdiffusion mechanism.526 As a result, the interfacial void are often reduced to a large extent, mitigating the sieve-in-a-cage morphology and increasing the membrane selectivity.527 A different approach to modify the filler surface is demonstrated by Shu and co-workers.528−530 Instead of using polymers, the modification was made by means of a Grignard reagent to create an inorganic nanostructured morphology such as nanowhiskers or asperities on the surface of the fillers. Generally, the mechanism involves a two-step reaction: (1) a dealumination reaction by thionyl chloride (SOCl2) to give sodium chloride (NaCl) and aluminum chloride (AlCl3) nanocrystals and (2) a succeeding reaction with methylmagnesium bromide (Grignard reagent, CH3MgBr) to grow Mg(OH)2 nanowhiskers on the filler surface using NaCl and AlCl3 as nucleating seeds (Figure 35a).530 The presence of Mg(OH)2 nanowhiskers creates a surface roughening effect which allows the polymer chains to intertwine onto the nanostructures in their randomly coiled conformations without the need for them to straighten into more orderly arrangements. In other words, adsorption of the polymer matrix onto the heterogeneous surface of a modified filler is entropically

Figure 35. (a) Scheme showing the synthesis of Mg(OH) 2 nanowhiskers on the surface zeolite 4A and (b) the change in interfacial morphology of the composite membranes before and after treatment. Adapted from ref 528. Copyright 2007 American Chemical Society.

driven and shows a more favorable Gibbs free energy as compared to the relatively smoother surface of an unmodified filler. This helps to promote a good polymer−filler adhesion as evidenced by the absence of interfacial voids in the SEM images of Figure 35b.528 With hindsight, the use of moisture-sensitive SOCl2 and CH3MgBr makes the overall synthesis of Mg(OH)2 nanowhiskers more complex and technically demanding. Besides, the reaction is typically limited to aluminosilicate fillers (typically zeolites) since NaCl and AlCl3 nanocrystals are indispensable nucleating seeds for the Mg(OH)2 nanostructures to grow.531,532 For this reason, the formation of Mg(OH)2 nanowhiskers on pure silica is not feasible and external sources via “surface seeding” (i.e., the addition of a small amount of zeolite crystals into the pure silica) is needed to ensure a successful synthesis of the Mg(OH)2 nanostructures on the pure silica surface.532 New methods have also been adopted to develop these nanostructures using simpler synthetic routes. For example, a facile solvothermal method was employed to deposit the inorganic nanostructures on a filler surface using magnesium sulfate (MgSO4) in a solution with a high ethylenediamine (EDA)-to-water ratio.520 When applying this method, the choice of magnesium source is critical as it was found that a substitution of MgSO4 for magnesium nitrate (MgNO3) resulted in Mg(OH)2 platelets rather than the desired nanowhiskers morphology. Although the exact cause has yet to be comprehended, it is speculated that the selective interaction of EDA with the surface Mg2+ ions impedes growth along specific lattice planes and discriminately limits the growth of Mg(OH)2 to only the 1D direction.520 Alternatively, it is possible to replace EDA with diethylenetriamine (DETA) in a modified solvothermal method owing to the corrosive nature of EDA and its strong AS


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the filler surface. As a result, this prevents the nonsolvent molecules from acting as nucleating agents, which inhibits the formation of a polymer lean phase around the filler and suppresses the development of defective interfacial voids.534 To have a better understanding of the underlying principle, detailed discussions are found in earlier reviews (see refs 480 and 481). 3.4.2.4. Room-Temperature Ionic Liquids. Room-temperature ionic liquids (RTILs) are organic salts that are in the molten phase at ambient conditions (Figure 37).536,537 Recently, RTILs have garnered the interest of the membrane community because of their many unique properties such as nonvolatility, high thermal stability, tunable chemistry, and good intrinsic solubility for CO2 gas.538,539 Moreover, they show good compatibility with an extensive range of inorganic and organic materials and can often bring materials of different nature into the same phase.540 Thus, RTILs have been widely employed for separation membrane or composite membrane preparations.135 To stay aligned with the scope of this review, we avoid venturing into a discussion of RTIL-based membranes (see ref 541 for a recent review on this topic) but rather briefly highlight the use of RTILs as a third component or agent for wetting of the polymer−filler interface in the composite membranes. The competitive advantages of using RTILs for composite membranes include: (1) stable dispersion of inorganic fillers (such as silica and metal nanoparticles) by the strong electrostatic repulsions of the RTILs on the functionalized filler surface,542 (2) tailorable gas transport performance by engineering CO2 solubility to create easy access to the functionalized filler and consequently enhancing the CO2/ CH4 solubility selectivity,543,544 and (3) effective mitigation of the sieve-in-a-cage morphology by having the RTILs to fulfill the role of a “wetting agent” for improving the polymer−filler interfacial adhesion. To this end, RTIL functionalization has been performed on NH2-MIL-101(Cr),545 ZIF-8,546,547 HKUST-1,548 and other fillers.549,550 In general, the RTILto-filler ratio, filler loadings, and the molecular design of the RTILs are found to be critical parameters for the filler-RTILpolymer interactions and performance of the composite membranes. Unfortunately, these parameters have yet to be systematically comprehended. Hence, combining an effective modeling of the three-component system (filler-RTIL-polymer) coupled with empirical observations can serve as a powerful strategy to close this gap. A recent study by Mannan et al. successfully demonstrated a modified Maxwell model,

tendency to plug the pores of the filler.520,533 On top of that, a method based on ion-exchange of Na+ framework cations with Mg2+ cations in the zeolite materials was also reported.533 Following the ion-exchange step, a hydrothermal reaction was carried out at a slightly basic medium (pH ∼ 9.5−11) to induce and precipitate the Mg(OH)2 nanowhiskers on the filler surface. As compared to the fillers modified by other methods, this method is clearly more environmentally friendly and produces superior modified-fillers that can increase the CO2/CH4 selectivity of the composite membranes without compromising the CO2 permeability (Figure 36).533

Figure 36. CO2/CH4 separation performances of unfilled polymeric (Matrimid) and composite membranes. The fillers involved are bare zeolite (Bare LTA) and surface-functionalized zeolite via various synthetic routes. Reprinted from ref 533. Copyright 2012 American Chemical Society.

Beyond all these techniques are other surface modifications which encompass methods for enhancing the surface hydrophobicity of the filler. This is usually realized either by chemical functionalization534 or deposition of a hydrophobic material535 on the surface of the filler. The objective is to mitigate the sieve-in-a-cage morphology, but, unlike the aforementioned techniques, increasing the surface hydrophobicity is an effective strategy only when the composite membrane is prepared via a nonsolvent phase inversion process. This is because the hydrophobic modifications can minimize nonsolvent (i.e., in most cases water) interaction at

Figure 37. Cations and anions of RTILs that are commonly adopted for CO2 separation. Reproduced with permission from ref 552. Copyright 2016 Elsevier. AT


zeolites, silica, MWCNTs, and mesoporous materials APTMS, APTES, and APDMES

well-established technique and increased resistance toward humid condition

additional defects such as rigidified interface and pore blockage can arise, gas permeability can be considerably sacrificed without improvement in selectivity, and careful optimization of loading and precise matching of polymer/filler is required

target filler

merit

limitation

example

covalent bonding between filler and polymer

silane coupling agent

principle

parameter

antiplasticization effect leading to huge decrease in gas permeability, limited choice of LMWMs owing to the need for optimal molecular weight, and LMWMs overdose can result in detrimental effect to the separation performance of the composite membranes

easier to administer than silane coupling agents and increased membrane selectivity

TAP; PNA, TBBPA, BHT, and HFBPA; and β-cyclodextrin

zeolites and MWCNTs

intermolecular forces of attraction between filler and polymer

low molecular weight material (LMWM)

only small amount of the polymers is needed, priming is useful for fillers that possess severe agglomeration (e.g., nanocrystals), and easier to administer than silane coupling agents inappropriate choice of nonsolvent during priming may lead to conglomeration and the role of sizing agents has yet to be fully elucidated, increasing the degree of uncertainty in the sizing technique

PVAc and PVP

compatible interaction of the dope polymer (priming) on the surface of the filler with the bulk polymer matrix and increased reactive sites or surface areas by the sizing agent on the surface of the filler for adhesion with polymer matrix zeolites and CMSs

priming/sizing agent

complex modification, technically challenging using moisture-sensitive chemicals like SOCl2 and CH3MgBr and corrosive amines like EDA, “surface seeding” necessary for silica due to absence of the Al or Na source, and choice of Mg source is critical for controlling Mg(OH)2 morphology

precursors include: grignard reagents (CH3MgBr), MgSO4 in EDA or DETA, and MgSO4 followed by NaNO3 (ion-exchange) at high pH (9.5−11) display strong attraction between filler and polymer and high selectivity obtainable without sacrificing gas permeability

zeolites and silica

inorganic Mg(OH)2 nanostructures on the surface of the filler to induce thermodynamically spontaneous polymer chains entanglement of the matrix

Mg(OH)2 nanowhisker

surface modification technique

modeling three-component permeation behavior (polymer-filler-RTIL) in composite membrane can be technically challenging and complex, lack of cost-effective RTILs with diverse functionalities for composite membranes, and poor scalability potential to support commercial-scale membrane fabrication

high stability in composite membrane owing to its nonvolatility and high thermal stability, RTILs can induce effective dispersion of fillers, and tailorable gas transport properties by increasing CO2/CH4 solubility selectivity

zeolites, MOFs, GO, MWCNTs, and metal oxide nanoparticles imidazolium-based RTIL and [Tf2N]-based RTIL

RTIL serves as a wetting agent to bring together the solid (filler) and polymer phase

room-temperature ionic liquid (RTIL)

Table 21. Overview of the Various Surface Modification Techniques and Representative Compatibilizers for Mitigating Non-Ideal Interfacial Morphologies in Composite Membranes


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Figure 38. Schematic illustration of the transport mechanism of CO2 and CH4 in a composite membrane with conventional zeolite fillers. Zeolite picture adapted with permission from ref 161. Copyright 2013 Royal Society of Chemistry.

not infallible. In some cases, interfacial defects are not fully restored and, as mentioned earlier, there remains the concern with the thermal stability of certain fillers when heating the composite membranes at a high temperature. From the practical point of view, there is no single ideal solution that can mitigate all types of nonideal morphologies in composite membranes. Therefore, it is important to take into account the physicochemical properties of the polymer matrices and fillers, as well as the techniques and conditions used in the fabrication of composite membranes when devising the most appropriate strategy for resolving nonideal interfacial morphologies. In this respect, an overview of the various surface modification techniques and representative compatibilizers are provided in Table 21. Along with the methods used in maintaining polymer chains flexibilities, this table serves as a quick guide to support readers in making informed choices about the different types of strategies that are available for mitigating nonideal interfacial morphologies.

which can accurately predict the CO2 permeability of a composite membrane comprising RTILs.551 Moving forward, similar modeling efforts can be put in place to identify the physicochemical properties of RTILs that are essential for improving interfacial defects and separation performances of the composite membranes. These efforts can also add value by providing a framework for guiding the design of novel RTILs to achieve the same purpose. 3.4.2.5. Maintaining Polymer Chains Flexibilities. The strategies discussed so far are focused on alleviating the interfacial incompatibility issues either by surface modification of the fillers or adding compatibilizers to strengthen the affinity and interaction between the polymer matrices and fillers (Table 21). Maintaining polymer chain flexibility is a unique strategy which uses a different approach to mitigate nonideal morphologies.481 The objective is to allow the polymer matrix to remain in the rubbery state for a long-enough period of time so that the polymer chains possess consistent mobility to diffuse freely and undergo sufficient relaxation in response to a stress. In this way, the polymer chains are sufficiently robust to deal with the uneven distribution of stress or inhibited contraction of the polymer that is usually associated with the incorporation of filler. Every method for maintaining the polymer chain flexibility has its own merits and limitations. The most straightforward method is to cast a polymer dope solution at a temperature higher than the Tg of the polymer matrix. Nevertheless, this requires the use of a low volatile solvent to avoid evaporation under the high temperature used during membrane casting.479 Many of the mainstream solvents for dissolving the polymer matrix are unable to meet this requirement, and this poses a big challenge for the successful implementation of this method. Alternatively, a polymer with a lower Tg can be chosen as the matrix but this drastically decreases the choice of polymer matrix, rendering the method equally impractical.116,553 In addition, the preparation of composite membranes using melt processing and the incorporation of a plasticizer into the polymer dope solution to decrease the Tg of the polymer have also been reported. However, the former method has little industrial relevance while the latter can result in compromised membrane performance since the plasticizers are known to induce a lower intrinsic separation performance of the polymer matrix.511,554 At present, the most widely used method is to carry out a postfabrication annealing of composite membranes at temperatures above the Tg of the polymer matrices.503 By doing so, the vitrified polymer chains regain their flexibility and this maximizes the capacity of the matrix to relax out the stress and heal the nonideal morphologies that are present in the composite membranes. Despite this, the annealing method is

4. PERFORMANCES OF COMPOSITE MEMBRANES There are several reasons why high CO2 uptake and CO2/CH4 selectivity of the fillers do not always lead to superior CO2/ CH4 separation performances of the composite membranes. One of the most prominent reasons is that the polymer (continuous) phase remains largely dominant in comparison to the solid (filler) phase of the composite membranes. As such, the permselectivity is expected to be primarily driven by the polymer phase of the composite membranes. In addition, strong interactions between the polymer and solid phase plays a significant role in producing defect-free composite membranes, which would otherwise result in compromised separation performances as already discussed in section 3. Hence, a critical analysis of CO2/CH4 composite membrane performances is necessary for understanding the effect and affinity of the fillers on the polymer phases. Notably, a synergistic effect, where both CO2 permeability and membrane selectivity are enhanced due to the incorporation of the filler, is most desirable. Therefore, in this section, we aim to explore the different types of functionalized materials and the chemistries involved, as well as the outcomes they create as fillers in composite membranes for biogas separation. 4.1. Conventional Materials

The gas transport mechanism across the composite membranes generally does not deviate from the conventional solution-diffusion mechanism. Ideally, the role of conventional fillers such as zeolites is to increase the CO2 adsorption capacity by providing additional reactive sites for interactions with CO2 molecules (Figure 38). This enhances the CO2 solubility of the composite membranes. Moreover, for fillers AV


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Figure 39. 40 wt % Na-LTA zeolite in PI matrix showing (a) low and (b) high magnification of the cross-sectional membrane morphologies. Reprinted with permission from ref 559. Copyright 2012 Wiley-VCH. (c) A comparison of the zeolite beta-PSf interfacial morphology before and after p-xylenediamine treatment. Reprinted with permission from ref 562. Copyright 2006 Wiley-VCH. (d) A plausible interaction between the zeolite beta and PSf matrix compatibilized via p-xylenediamine. Reproduced with permission from ref 562. Copyright 2006 Wiley-VCH.

with suitable pore windows, the smaller CO2 molecules are also able to access the pores and acquire a direct diffusion pathway through the filler and across the composite membrane. On the other hand, the larger CH4 molecules typically experience less favorable interactions and are unable to access the pores of the filler. Thus, the CH4 solubility is relatively lower and the molecules assume a more tortuous diffusion pathway through the composite membrane, which decreases the CH4 diffusivity (Figure 38). Correspondingly, the separation performance of composite membranes demonstrates a higher CO2 permeability and, in the case of an ideal interfacial morphology, a greater distinction between the CO2 and CH4 permeability also leads to an increase in the CO2/ CH4 membrane selectivity. In this regard, the following sections are focused on the effect of various functionalized materials in composite membranes and how they can be leveraged to achieve synergistic performance enhancements for transcending the Robeson upper bound limit. 4.1.1. Zeolites and Related Materials. Zeolites are one of the most common fillers for composite membranes given their compelling properties such as attractive CO2 adsorption, size/shape selectivity through their well-defined pores, as well as high chemical and thermal stability, which allow greater versatility in the membrane fabrication process to yield composite membranes with minimal defects. The first reported zeolite composite membrane was made in 1973 when Paul and Kemp555 demonstrated a CO2 diffusion time-lag by embedding microporous zeolite 5A within the PDMS polymer matrix. Following this, there are many other studies in the literature, which pointed out that the gas transport properties of zeolitebased composite membranes are strongly influenced by both the zeolites and polymers. Hence, judicious matching of the correct type of zeolites to the right kind of polymer matrix is

considered a key parameter for the success of the zeolite-based composite membranes. Initial studies on zeolite-based composite membranes were mainly focused on the use of LTA zeolite as filler.429 For instance, Bahtiyar and Firuze utilized LTA zeolite with different pore sizes such as K-LTA (3 Å), Na-LTA (4 Å), and Ca-LTA (5 Å) to fabricate composite membranes for CO2/CH4 separation.118 It was found that an increase in the pore size of the LTA zeolite from K- < Na- < Ca-LTA resulted in an increase in the CH4 permeability with a corresponding decrease in the overall CO2/CH4 selectivity. Besides, the study also highlighted the importance of the choice of polymer by incorporating zeolites of different loadings into the PEIm and PI polymer matrices. Clear synergistic enhancements in both CO2 permeability and CO2/CH4 selectivity were observed when the PI membrane was loaded with 30 wt % zeolite CaLTA, pushing the performance close to the 2008 Robeson upper bound limit.118 On the contrary, the PEIm composite membrane failed to show this synergistic effect, owing to a poorer zeolite-polymer interfacial adhesion. In another study, Li and co-workers investigated the use of larger electropositive cations like Ag+ for ion-exchanged LTA zeolite.119 The work was motivated by the exploitation of larger ionic radius of Ag+ (1.15 Å) to drive a further reduction in the pore size of the AgLTA zeolite. As a result, the pore size of Ag-LTA was reduced close to the kinetic diameter of CO2 and this caused a strong decline in the CH4 permeability of the composite membrane. Furthermore, facilitated transport via π-bonded complexes between Ag+ and CO2 also led to a stronger CO2 adsorption as compared to its Na-LTA counterpart. Such a facilitated transport mechanism has been reported in many composite membranes in the literature (see section 1.3.1).556−558 In a similar fashion, other zeolite categories are adopted in composite membranes. For example, Karkhanechi et al. AW


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tend to result in a significant loss in the CO2 permeability (see section 3.4.2.2). The development of hierarchical zeolites is another approach which shows great promise in reducing the likelihood of nonselective interfacial void between the zeolite and polymer matrix. To synthesize hierarchically structured zeolites, cetyltrimethylammonium bromide (CTAB) is often employed as a template for the mesoporous material synthesis. Particularly, zeolite ZSM-5 with mesopores of ∼2.7 nm was synthesized and characterized based on a Barrett−Joyner− Halenda (BJH) desorption profile.565 Zhang and co-workers incorporated a 10 wt % loading of this mesoporous filler into Matrimid and found an improved molecular sieving effect given that the polymer matrix was able to penetrate and constrict the mesopores of the filler. More importantly, due to this penetration, interfacial adhesion between the zeolite and polymer matrix was enhanced, leading to a 93.6% increment in the CO2/CH4 selectivity. Despite these promising results, commercial acceptance of zeolite-based composite membranes is hindered by poor mechanical properties,565,566 especially for zeolite-based TFC membranes and high zeolite loadings, which results in poor dispersion within the polymer matrix.567,568 4.1.2. Metal−Organic Frameworks. MOFs exhibit many attractive properties such as high surface areas, low densities, tailorable porosities, and tunable physicochemical properties, making them suitable as fillers for composite membranes. Moreover, the organic moieties (ligands) of MOFs are often attributed to the improvement in the affinity between the crystalline MOF and organic polymer matrix (see section 2.1.2). Hence, MOFs possess a greater capacity than zeolites to mitigate the sieve-in-a-cage morphology that may be present in composite membranes. In 2004, the first MOF-based composite membrane was reported where 20 and 30 wt % of copper(II) biphenyl dicarboxylate-triethylenediamine were loaded into a poly(3-acetoxyethylthiophene) (PAET) matrix to yield composite membranes with a noticeable increase in the CH4 permeability.112 This result was an outcome of an enhanced membrane hydrophobicity and CH4 affinity brought about by the incorporation of the MOF filler. Since then, the approach of composite membranes using MOFs as fillers has been extensively explored for CO2/CH4 separation. To date, MOFs which include HKUST-1, MOF-5, CuBDC, ZIF-7, ZIF8, ZIF-90, MIL-53, MIL-101, SIFSIX-3-Zn, and UiO-66 have been incorporated into composite membranes.18,149,569−572 Rubbery as well as glassy polymers such as PDMS, Pebax, Matrimid, PSf, Ultem, and 6FDA-based PI have also been widely utilized as a continuous phase of these composite membranes (Table 22 to Table 27). Controlling the particle size of MOFs is one of the most critical parameters for the successful fabrication of MOF-based composite membranes. As demonstrated by Gong and coworkers,106,571 sedimentation of fillers at the bottom of flat sheet composite membranes was observed for bulk SIFSIX-3Zn crystals but not when the SIFSIX-3-Zn crystals were in the submicron- and nanosized range (Figure 40). Accordingly, a more uniform distribution of smaller-sized SIFSIX-3-Zn crystals in a XLPEO matrix was obtained at 10 wt % loading. This gave a CO2 permeability of 620 barrer and CO2/CH4 selectivity of 27, which corresponded to a synergistic enhancement of 38 and 80% in both the CO2 permeability and CO2/CH4 selectivity, respectively. Furthermore, when the smaller-sized SIFSIX-3-Zn crystals were matched with three different 6FDA-based PI, the composite membranes were

synthesized Na-FAU and Na-LTA zeolites and incorporated them into a PI polymer matrix to fabricate composite membranes.559 The resulting composite membranes were obtained without appreciable interfacial defects, even at a high filler loading of 40 wt % (Figure 39a,b). The best membrane comprising 30 wt % loading of Na-FAU zeolite exhibited a synergistic increase in both the permeability and selectivity (96 and 280%, respectively) when compared to the unfilled PI membrane. This increase was, however, not found for Na-LTA zeolite because of its smaller pore size that led to a slow diffusion of permeating gases. Additionally, the potential of chabazite zeolites SAPO-34 and SAPO-44 for CO2/CH4 separation were also evaluated.560,561 In these studies, a PSf polymer matrix was chosen to prepare asymmetric composite membranes. As the mass transfer resistance was lower, given its asymmetric morphology, the CO2 permeance demonstrated a 4- to 10-fold increment (from 22 to 350 GPU) as compared to an unfilled PSf membrane. Enhancements in CO2/CH4 selectivity were also achievable at loadings of 10 and 5 wt % for SAPO-34 and SAPO-44, respectively. Beyond these optimized loadings, the sieve-in-a-cage morphology became dominant, leading to less selective separations. Driven by the numerous advantages of the hollow fiber membrane configuration (see section 1.3.2), zeolite-based hollow fiber membranes were also of interest for potential industrial use. Specifically, dual-layer hollow fiber composite membranes were prepared by utilizing Matrimid as the inner support layer, while zeolite beta particles were incorporated into the selective ultrathin PSf outer layer.562 Subsequently, the as-spun composite fibers were treated with a p-xyelenediamine compatibilizer to afford strong adhesion between the zeolite particles and PSf polymer matrix (Figure 39c,d). This contributed to a substantial improvement in the CO2/CH4 selectivity from 24 to 39 at a filler loading of 30 wt %, while the CO2 permeance remained uncompromised (from 2.5 to 1.5 GPU).562 To better enhance the adhesion between the zeolite and polymer matrix, Hudiono et al.542,563 presented a threecomponent composite membrane comprising RTIL as a compatibilizer, zeolite as a filler, and the polymer matrix. Although such an approach has been utilized by the Universal Oil Products (UOP)564 on ACs, its application on zeolitebased membranes was only recently introduced. Typically, RTILs serve as a “wetting agent” to bring together two uniquely different phases such as the inorganic zeolite and organic polymer. This helps to resolve the organic−inorganic incompatibility issue and induce a strong adhesion for mitigating the sieve-in-a-cage interfacial morphology between the zeolite and polymer matrix (see section 3.4.2.4). Besides, RTILs also possess a high intrinsic solubility for CO2 gas. By virtue of this, the CO2 permeability of a SAPO-34-[C2mim][Tf2N]-poly(RTIL) composite membrane exhibited a 418% enhancement in permeability without substantial compromise in the CO2/CH4 selectivity as compared to the reference composite membrane without RTIL.542 In addition, by increasing the amount of [emim][Tf2N] in the composite membrane from 12 to 36 wt %, the membrane with 40 wt % SAPO-34 loading showed a 753% increase in the CO2 permeability without loss of CO2/CH4 selectivity.563 Hence, the use of RTILs to mitigate poor filler adhesion in zeolitebased composite membranes shows a clear competitive advantage in contrast to other strategies like LMWMs, which AX


AY

metal oxide nanoparticles

silicas

CNTs

layered silicates

GFMs

carbon-based particles

MOPs

ZIFs

MOFs

zeolites

type Pebax-1657 XLPEO Pebax-1657 Pebax-1657 PDMS PVAc PDMS Pebax-1657 Pebax-1657 Pebax-2533 PVC-g-POEM PVC-g-POEM Pebax-1657 Pebax-1657 Pebax-1657 ABS ABS PDMS PDMS PDMS Pebax 1657 Pebax 1657 Pebax 1657 Pebax 1657 PDMS PDMS Pebax 1657 Pebax 1657 EVA Pebax-2533 Pebax-1657 Pebax-1657 Pebax-1657 PB PB

HKUST-1 HKUST-1 NH2-HKUST-1 ZIF-8 ZIF-8 ZIF-8 (hollow) IL@ZIF-8 PDA [emim][BF4] AC-1

AC-2 nonfunctionalized graphene GO GO GO PEI-PEG-GO PEI-PEG-GO Im-GO TMAB-MMT sepiolite NIPAM-CNT PEG-MWCNT fumed silica octanol-functionalized Si nanoparticle Ag

Ag/IL Al2O3 + PEG1000

MgO brookite TiO2

rubbery matrix

Na-A amine Ca-A SAPO-34 ZSM-5 ZSM-5 CuTPA

filler

Ag: 0.005; IL: 0.5 Al2O3: 8; PEG1000:10 27 27

0.005

40 0.5 1 8 1 10 10 0.8 15 20 5 3 10 33.3

30 20 20 35 30 30 15 5 80 10

5 10 50 10 66 15

filler loading (wt %)

15 15

10 25

10

2 1.1 10 10.1 3 2 2 4 13.6 13.6 2 2 8 2

4.9 1 7 1 20 CO2: 0.09 CH4: 4.48 − 3 3 2 0.1 1 1 1 3 2

pressure (bar)

35 35

35 25

35

560 230

180 71.6

220

20.50 4460 2429.76 142.22 100.00 145.00 1330.00 64.00 − − 580 35.00 46 400

2900 167.3 170.1 1287 687.7 623.0 104.8 69 270 10.81

− 30 30 25 35 35 25 30 35 20 20 37 35 35 25 30 30 25 − − 25 25 25 25

71.4 380 338 192.7 11648 3.26

P(CO2) (barrer)

25 40 35 35 27 35

temp (°C)

testing condition

1300.00 379.17

63.6 −

100

614.3 47.68 −40.69 −60.00 102.06 76.83 170.87 3.56 − − 93.33 64.47 109.09 32.89

16 98.7 102.0 267 880 790 45 13.1 212.5 131.2

28.0 20.0 204.5 57.4 232.9 34

permeability enhancement (%)a

5.09 10.95

61.02 24.1

36.67

50.5 4.2 3.58 9.40 24.66 24.0 45.0 25.10 11 14 35 23.78 9.89 9.40

3.6 19.5 26.2 9.0 12.4 11.2 34.9 18 27.3 34.3

32.2 18 17.8 32.66 4.36 40.4

α(CO2/CH4)

separation performance

3.34 − 1.91 2.17

−24.02 25.86

2.19

3.92 0.80 −0.10 2.07 1.06 1.19 3.27 −0.07 − − 0.90 −0.14 1.74 0.28

0.54 1.14 1.94 1.51 1.96 1.66 2.09 0.60 2.50 1.77

1.78 −0.35 1.22 1.61 1.98 0.68

Findex value

194.1 −

76.7

109.5 16.67 17.6 209.90 14.42 26.32 136.84 −3.83 83.33 133.33 9.38 −21.36 46.09 0.00

16.1 18.9 59.8 8.4 −11.4 −18 92 20 67.48 42.3

78.9 −18.4 4.0 55.0 34.2 15.8

selectivity enhancement (%)

696 113

694 695

694

594 117 687 688 610 612 612 689 690 690 691 692 111 693

682 683 683 102 684 105 547 685 686 594

678 679 680 101 681 252

ref

Table 22. Enhancements in the Pure CO2 Permeability and CO2/CH4 Selectivity of Symmetric Composite Membranes by Using Different Fillers in Polymer Matrices



Review

predicted by the Maxwell model to be capable of achieving separation performances that were beyond the upper bound limit.106 This was later experimentally demonstrated using a composite membrane comprising 20 wt % loading of SIFSIX3-Zn nanocrystals in a 6FDA-TMPDA matrix. Indeed, a high CO2 permeability of 1120 barrer and CO2/CH4 selectivity of 38 was achieved as a result of reducing the particle size of the filler. Essentially, these results suggest that smaller MOF crystals have an absolute advantage in terms of their stabilities in dope solutions because the polymer chains can wrap the fillers more effectively. This not only ensures a better dispersion of the MOF fillers in the polymer matrix but also offers stronger polymer−filler interactions.571 Enhancing this interaction can also be instrumental for increasing the resistance of MOF-based composite membranes toward CO2-induced plasticization.266 Apart from the particle size, the morphological structure of MOF fillers can be leveraged to engineer the separation performances of composite membranes. Specifically, 2D MOFs are of great interest given their promise as effective barriers for composite membranes. They exhibit high aspect ratio, nanoscale thickness, and cavities of suitable pore windows that only allow smaller CO2 molecules to pass through but exclude larger CH4 molecules. As a result, the impassable CH4 molecules have to adopt a more tortuous diffusion pathway around the 2D MOFs, leading to a larger decrease in the CH4 diffusivity and an enhanced CO2/CH4 diffusivity selectivity (Figure 41a).573 For these reasons, composite membranes comprising 2D MOFs often achieved improved CO2/CH4 separation at a lower filler loading as compared to the other fillers with different morphologies. A low filler loading has the advantages of less possibility of filler agglomeration, better preservation of the mechanical properties, and potential reduction of the composite membrane thickness. Hence, the strategic use of 2D MOFs can provide solutions to many of the challenges faced by composite membranes.149 As such, Rodenas et al.251 capitalized on a bottom-up approach to synthesize nanosheets of copper 1,4-benzenedicarboxylate (nsCuBDC) with micrometer lateral dimension and nanometer thickness (Figure 41b,c). In this study, the incorporation of nsCuBDC into a Matrimid polymer endowed the membranes with CO2/CH4 selectivities that were up to 8-fold higher than its bulk CuBDC counterpart, and 30 to 80% larger as compared to the unfilled Matrimid membrane. Nevertheless, the CO2 permeability was compromised to about 20 and 30%, respectively.251 In place of the Matrimid matrix, Yang and coworkers573 recently integrated ns-CuBDC into rigid polymer matrices of high free volumes. With a low ns-CuBDC loading of just 2 and 4 wt %, composite membranes using PIM-1 and 6FDA-DAM matrices were able to demonstrate up to 40% increase in the CO2/CH4 selectivity. The incentive of this approach was to capitalize on the high intrinsic CO 2 permeability of the polymer matrix to offset the decrease in the CO2 permeability brought about by the incorporation of the 2D MOFs. Hence, despite a 25 to 35% decline in the CO2 permeability, the separation performances were able to surpass the Robeson upper bound limit given their enhancements in CO2/CH4 selectivity.573 Critically, this approach strengthens the competitive position of 2D MOFs as advanced fillers for high-performance composite membranes. In this regard, adopting a suitable polymer matrix to match the nature of the MOF fillers is a pragmatic approach to design MOF-based membranes. This strategy not only works for

a Permeability enhancement is the difference in the CO2 permeability between the composite and unfilled membranes expressed as a percentage of the CO2 permeability of the unfilled membrane. This is equivalent to the facilitation ratio, which is an index previously defined by Jia and co-workers697 to benchmark performances of zeolite-based membranes (see section S4.1).

Table 22. continued

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Figure 40. Cross-sectional morphologies of the composite membranes using (a) bulk and (b) submicron SIFSIX-3-Zn crystals and FESEM images showing the particle size of the (c) bulk and (d) submicron crystals. Reprinted with permission from ref 106. Copyright 2015 Elsevier.

hydroxyl groups facilitated hydrogen bonding with the −NH2 groups of the NH2-MIL-53(Al) fillers, leading to better adhesion between the MOF and polymer matrix. In comparison, the nonfunctionalized MIL-53 fillers were unable to exert the same interaction owing to the lack of −NH2 groups. This brought about a decrease in the CO2/CH4 selectivity for the 6FDA-(DAM)-(HAB) membrane loaded with 10 wt % MIL-53(Al) while its NH2-MIL-53(Al) counterpart demonstrated a remarkable 4.4-fold increase in selectivity as compared to the unfilled membrane.73 Considering the ability to form hydrogen bonding with the polymer chains, MOFs containing amine functional groups are deemed compelling for mitigating the polymer−filler interfacial nonidealities. Apart from the above example of 6FDA(DAM)x-(HAB)y, the NH2-MIL-53(Al) filler has also shown excellent adhesion with different polymer matrices such as PVDF and PSf.127,574,575 For instance, Zornoza et al.576 reported the presence of hydrogen bonding between the amine moiety of NH2-MIL-53(Al) and the sulfone groups of the PSf chains. Due to this favorable interaction, a filler loading, as high as 40 wt %, was made possible. The best performing membrane was, however, obtained with a 25 wt % loading of NH2-MIL-53(Al), which exhibited a near 2-fold increase in the CO2/CH4 selectivity relative to the unfilled PSf membrane (46 vs 23). Notably, at a high pressure of 10 bar, the CO2/CH4 selectivity was further enhanced to 110 while achieving a higher CO2 permeability than the reference membrane. This result, at a high transmembrane pressure, also showed great promise in using NH2-MIL-53(Al) as a filler to overcome the plasticization effect common in polymeric membranes (see section 1.3.3). The enhancement in selectivity was attributed to the “breathing” effect from the intrinsic flexibility of the MOF filler, which expanded at high pressure to allow additional CO2 loading in the pores (see section 2.1.2.2). Such an expansion was believed to close polymer−filler interfacial voids leading to enhanced gas selectivity, while the

Figure 41. (a) Schematic illustration showing the role of 2D MOFs in a composite membrane. Reprinted with permission from ref 573. Copyright 2017 Royal Society of Chemistry. FESEM images showing (b) ns-CuBDC and (c) its respective membrane surface morphology. Reprinted with permission from ref 251. Copyright 2015 Nature Publishing Group.

MOF fillers of different morphological structures but also for MOF fillers with different chemical functionalization. As reported by Tien-Binh and co-workers,73 a series of MIL53(Al)- and NH2-MIL-53(Al)-based composite membranes were prepared from three different 6FDA-based coPIs (i.e., 6FDA-(DAM)x-(HAB)y with a DAM to HAB mole ratio of 2:1, 1:1, and 1:2). It was found that a greater substitution of DAM with HAB gave a stronger interpolymer chain interaction that was induced by the hydroxyl groups of the HAB copolymer. This was accompanied by a decrease in the CO2 permeability but an enhancement in the CO2/CH4 selectivity of the purely polymeric membranes. More importantly, the BA


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Figure 42. Schematic illustration showing the (a) effect of ILs on the reduction of the effective nanocage size. Reprinted with permission from ref 584. Copyright 2015 Wiley-VCH. (b) Incorporation of polydopamine (PD) onto the surface of ZIF-8 nanocrystals with FESEM images of composite membranes containing (c) 20 wt % ZIF-8 and (d) 20 wt % ZIF-8@PD in PI matrix. Reprinted with permission from ref 585. Copyright 2016 Wiley-VCH. (e) Cross-sectional membrane morphologies revealing good adhesions of ZIF-90 crystals in (1) ZIF-90A/Ultem, (2) ZIF-90A/ Matrimid, (3) ZIF-90A/6FDA-DAM, and (4) ZIF-90B/6FDA-DAM composite membranes. (f) Robeson plot showing the performance enhancements of the ZIF-90 composite membranes. Reprinted with permission from ref 586. Copyright 2010 Wiley-VCH.

greater CO2 loading within the pores of the filler contributed to the higher CO2 permeability of the composite membrane.576 On top of that, the flexible framework of NH2-MIL-53(Al) also appears to be dependent on the composite membrane casting conditions. As verified by Rodenas et al.,577 fast solvent evaporation initiated the contraction of the NH2-MIL-53(Al) filler in a Matrimid matrix. Hence, a higher percentage of the MOF took on a narrow pore configuration that prevented the penetration of polymer into the pores. This resulted in a 70% increase in the CO2 permeability, combined with a slight enhancement in CO2/CH4 selectivity for the composite membrane with a 25 wt % filler loading.577 Recently, Tien-Binh and co-workers572 employed a novel approach for composite membrane preparation where

functionalized UiO-66-NH2 fillers were introduced as crosslinkers to react with the monomers of PIM-1 via an in situ polymerization. As a result of this in situ cross-linking reaction, the MOF-based composite membrane with 20 wt % filler loading exhibited excellent polymer−filler adhesion, which led to synergistic enhancements of 90 and 159% in the CO2 permeability and ideal CO2/CH4 selectivity, respectively. Likewise, a synergistic enhancement in the separation performance was also observed during a 1:1 CO2/CH4 mixed-gas evaluation. More strikingly, the cross-linked MOF-based membrane reported a stable CO2/CH4 selectivity over the pressure range from 2 to 10 bar and was found to be less susceptible to physical aging as demonstrated by a mere 23% decrease in the CO2 permeability (compared to an 83% BB


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clusters of the fillers and the PA domains of Pebax, leading to compatible interactions at the polymer−filler interface and the toughening effect.547 In another study, the compatibility between the filler and polymer was improved by coating a thin and uniform layer of polydopamine (PD) over the surface of ZIF-8 nanocrystals (Figure 42b).585 The ZIF-8@PD-based composite membranes were then prepared using a Tröger’s base-based PI as the polymer matrix. Although the CO2 permeability of the membrane containing 30 wt % ZIF-8@ PD was slightly lower than its counterpart, which contained the unmodified ZIF-8, the CO2/CH4 selectivity was enhanced as a consequence of the PD coating. Explicitly, the CO2/CH4 selectivity increased from 16 to 20 while the CO2 permeability declined from 1437 to 1056 barrer when ZIF-8@PD was utilized as fillers. Analogous to ZIF-8, ZIF-90 has a sodalite cagelike structure and a 0.35 nm pore window. Hence, it is also exploited for composite membranes and shows similar promise as that of ZIF-8. Essentially, ZIF-90 possesses two main benefits over ZIF-8: (1) the presence of carbonyl groups in the imidazolate linkers of ZIF-90 enables favorable intermolecular interaction with the CO2 molecules and (2) the superior thermal stability as characterized by the unimpaired crystallinity of ZIF-90 at elevated temperatures. However, unlike ZIF-8 which displays a wide and customizable particle size, the size of ZIF-90 crystals synthesized via a conventional solvothermal method is normally fixed and generally too large (∼100 μm) for effective incorporation into composite membranes. The synthesis of submicron or nanosized ZIF-90 is also not commonly reported. Among the small number of reports, a study by Bae et al.586 successfully synthesized submicrometer-sized ZIF90 crystals via a nonsolvent-induced crystallization method and incorporated them into polymer matrices such as Matrimid, Ultem, and 6FDA-DAM. Cross-sectional membrane morphologies revealed excellent adhesions of the ZIF-90 filler with the various PIs, which evidently led to significant enhancements in their CO2 permeabilities (Figure 42e,f). Nonetheless, when ZIF-90 was matched with Matrimid and Ultem, which have CO2 permeabilities far lower than that of ZIF-90, insignificant changes in the CO2/CH4 selectivity of the membranes were observed. On the contrary, matching ZIF-90 with a more permeable 6FDA-DAM PI resulted in a substantial enhancement in both CO2/CH4 selectivity and CO2 permeability (Figure 42f). Once again, this emphasizes the importance of careful selection and matching of filler and polymer matrix for composite membrane fabrication. 4.1.4. Microporous Organic Polymers. In recent years, MOPs have attracted immense attention as excellent fillers for composite membranes in view of their high surface areas, lowdensities, high structural and chemical tunabilities, as well as relative ease of functionalization. More distinctively, MOPs are structurally organic, which means they can create a stronger chemical affinity with the polymer matrix and are more compatible as compared to the inorganic fillers discussed so far. Hence, MOPs display greater dispersions in the polymer dope solutions and tend to form membranes with better structural integrity, leading to enhancements in both the CO2 permeability and CO2/CH4 selectivity of the composite membranes. One good example is the Schiff-based SNW-1 (Figure 43a). It has intrinsic ultramicroporosity with pore dimension of 5 Å and abundant amine functional groups as active sites for CO2 capture. Owing to these reasons, SNW-1 has been utilized as a

decrease for the unfilled PIM-1 counterpart) over a 400-day period.572 Hence, this illustrates that MOF-based composite membranes have the potential to meet the challenge of physical aging and CO2-induced plasticization in purely polymeric membranes (see section 1.3.3). Despite these promising results, MOFs typically possess lower chemical, thermal, and hydrolytic stability relative to the other common inorganic fillers, which include zeolites and CMSs. Particularly, the hydrolytic instability of certain MOFs is a major limitation, in view of the significant amount of water vapor content in the as-produced biogas (Table 1).578 To this end, only ZIF-8, MIL-53, MIL-101, and UiO-66 are considered water-stable among the MOFs that are commonly used in CO2/CH4 separation membranes.579,580 This was clearly demonstrated in the studies by Xin et al.,129,581 where respective composite membranes of MIL-101(Cr) and its functionalized derivatives showed stable and even increased CO2/CH4 selectivities when subjected to humidified mixed-gas feeds. However, to further harness the potential of MOF-based composite membranes for CO2/CH4 separation, it is necessary to move beyond these few types of MOFs. Hence, we contend that future work should focus on two strategic areas to address this current limitation in MOF-based composite membranes: (1) designing next-generation water-stable MOFs with stronger metal−ligand bonds as well as greater capacities to mitigate water vapor accessibility through improved hydrophobicity and steric hindrance approaches and (2) realizing composite membranes of these MOFs with better understanding of their long-term stability under realistic humid conditions. 4.1.3. Zeolitic Imidazolate Frameworks. Owing to the limitations inherent in MOFs, ZIFs emerge as a more viable option given their exceptional chemical and thermal stabilities in addition to other desirable merits that are similarly exhibited by MOFs. In particular, ZIF-8 is a popular choice for composite membranes, and it has since been incorporated into many different polymer matrices such as PIM-1,582 6FDAdurene,583 and Matrimid275 for applications in various gas separations including paraffin/olefin separation. However, ZIF8-based composite membranes demonstrate a smaller CO2 separation potential because of the low CO2 uptake capacity and modest CO2/CH4 selectivity of ZIF-8 (Table 12).276 Lately, research efforts have been focused on improving these shortcomings of ZIF-8. In particular, the size of the nanocage of ZIF-8 can be restricted by using imidazoliumbased ionic liquids (ILs) such as [bmim][Tf2N].584 When [bmim][Tf2N] was confined within the nanocage of ZIF-8, the nanocage size effectively shrunk from 1.12 to 0.59 nm (Figure 42a). This resulted in a lower amount of CH4 adsorbed in the nanocage, and the CO2 uptake was enhanced as a result. The composite membranes derived from this IL@ZIF-8 filler and the PSf matrix hence exhibited permselectivity that transcended the CO2/CH4 upper bound limit. Moreover, the CO2/CH4 selectivity was improved with increasing pressure, suggesting a high relevance for natural gas purification.584 The IL@ZIF-8 filler can also create a toughening effect at the polymer−filler interface, which increases the mechanical properties of the composite membranes. Li et al. 547 demonstrated this by adding 15 wt % IL@ZIF-8 filler into a Pebax matrix to afford a composite membrane with 20 and 280% higher tensile strength and elongation at break, respectively, than the pure Pebax counterpart. They proposed a hydrophobic−hydrophobic interaction between the IL BC


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43c). As a result, a membrane containing PAF-1-Li6C60 exhibited an overall 50% reduction in the CH4 permeability over a 365-day span, while the CO2 permeability remained generally unaffected. In the light of such selective aging, the PAF-1-Li6C60 filler can provide a competitive edge through improving the CO 2/CH4 selectivity of the composite membrane with increasing aging duration.125 A family of water-stable COFs is also synthesized via a combination of reversible Schiff base reaction and irreversible enol-to-keto tautomerization process (Figure 43d).322,587 These materials display robust hydrolytic stabilities in boiling water and acidic solutions given the irreversible nature of the tautomerization process and the absence of imine bonds in the keto form of the COFs. Two COFs, namely, TpPa-1 and TpBD,322,587 were selected for composite membrane preparation using tertiary butyl-substituted PBI (PBI-Bul) as the polymer phase.588 A high loading of 50 wt % was achieved with both COF fillers due to the favorable hydrogen bonding between the amine groups of the filler and the benzimidazole groups of the PBI matrix. The CO2 permeability of both COFbased membranes was also linearly enhanced with an increasing loading from 20 to 50 wt %. Comparatively, at 50 wt % loading, the CO2 permeability of the membranes containing the TpBD and TpPa-1 filler was 14.8 and 13.1 barrer, respectively. This difference was a result of the slightly larger pore aperture of TpBD as compared to TpPa-1 (24 vs 18 Å) (Table 13). Despite some marginal decline in the CO2/ CH4 selectivity due to filler agglomeration, the CO2/CH4 separation performance of these COF-based membranes was much closer to the Robeson upper bound limit, demonstrating the impressive capabilities of well-ordered MOPs as filler materials.588 To further unlock the promise of MOP-based composite membranes, we propose the following two strategic areas of focus. First, designing MOPs with ultrahigh surface areas and tailoring the morphological structures, as well as pore and particle sizes, are crucial in maximizing the capacity of MOPs for membrane-based biogas upgrading. At present, there are only a few classes of MOPs with surface areas that can match that of MOFs (BET surface areas of >4000 m2 g−1). It is also generally harder to control precisely the pore size of most amorphous MOPs except for ordered crystalline COFs where pore sizes are defined crystallographically (see section 2.1.4).589,590 Furthermore, owing to challenges in realizing materials synthesis and colloidal stability,591,592 the effect of morphology and physical size of MOPs on membrane performances is rarely investigated and thus poorly understood hitherto.156,593 In this respect, we believe that greater efforts should be made to optimize the rational design of MOPs for CO2/CH4 separation membranes. Second, exploiting unique advantages of MOPs is technologically important for the success of MOP-based membranes. Particularly, the thermal and chemical stability, as well as synthetic versatility and diversity of MOPs, are potentially valuable for strengthening the robustness and stability of composite membranes.590 Moving forward, there is a critical need to systematically study and evaluate the mechanical stability and long-term performance of MOP-based membranes, especially under harsh separation conditions. 4.1.5. Carbon-Based Particles. 4.1.5.1. Activated Carbons. The first AC-based membrane was pioneered by Kulprathipanja and Charoenphol564 for the study of CO2/N2 separation. An increase in the CO2/N2 selectivity was reported

Figure 43. (a) The chemical structure and molecular model of a SNW-1 fragment and (b) the cross-sectional morphology of a SNW1/PSf membrane at 12 wt % loading. Adapted with permission from ref 321. Copyright 2014 Wiley-VCH. (c) Separation performance of membranes comprising various derivatives of PAF-1 with PTMSP tracked over 365 days of physical aging. Reprinted from ref 125. Copyright 2015 American Chemical Society. (d) Scheme showing the synthesis of TpPa-1 and -2 via a reversible Schiff base reaction and an irreversible enol-to-keto tautomerization. Reprinted from ref 587. Copyright 2012 American Chemical Society.

filler in PSf membranes prepared via a spin-coating method.321 The organic nature of SNW-1 enabled excellent adhesion of the nanoparticles with the PSf matrix as evidenced by the intimate polymer−filler interfacial gap as shown in Figure 43b. At 12 wt % filler loading, the composite membranes demonstrated higher CO2/CH4 mixed-gas selectivity than the ideal selectivity (34 vs 27) due to the competitive adsorption of CO2 over CH4, as well as the presence of the molecular sieving effect exerted by the small micropores of SNW-1. A gradual increase in the CO2/CH4 selectivity was also observed with increasing SNW-1 content until the filler loading reached an optimum value of 12 wt %, beyond which a decrease in the separation performance occurred due to the appearance of defects and large interfacial gaps. Furthermore, given the strong affinity of CO2 with the amine functional groups of SNW-1, the CO2 permeability exhibited a 120% enhancement at 22.4 barrer over the unfilled PSf membrane.321 To further evaluate the effect of chemical functionality of MOPs on the CO2 permeability, PAF-1 and its functionalized analogues, PAF-1-NH2, PAF-1-SO3H, PAF-1-C60, and PAF-1Li6C60, were exploited as fillers for composite membranes using PTMSP as the continuous phase.125 The CO2 permeability of all these composite membranes was found to improve by at least 50% relative to the unfilled PTMSP membrane. In particular, a composite membrane containing PAF-1-Li6C60 showed the highest increment of 70%, reaching a remarkable CO2 permeability of 50,600 barrer. This was attributed to the additional sorption sites from the Li atoms and the ultrahigh surface area generated by the C60 nanoparticles. Most notably, these fillers value-added the PTMSP matrix by imparting selective aging properties to the composite membranes (Figure BD


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Figure 44. Comparison of separation performances of composite membranes with 20 wt % filler loading using (a) dry and (b) wet CO2/CH4 mixed-gas ratio of 1:9. Reprinted with permission from ref 578. Copyright 2015 Elsevier.

Figure 45. Robeson plot showing the effect of (a) CMSs loading on the CO2/CH4 separation performances of Ultem and Matrimid membranes. Prepared with data from ref 599. (b) Effect of DEA addition on the separation performances of CMSs/PES membranes where M-1 to -5 represent pure PES, CMSs/PES, 5, 10, and 15 wt % DEA in CMSs/PES membrane, respectively. Reprinted with permission from ref 600. Copyright 2015 Royal Society of Chemistry.

AC-based membranes, in particular Carbon C, demonstrated relatively stable CO2/CH4 separation performances using a wet mixed-gas feed while their counterparts using HKUST-1, ZIF8, and UiO-66 as fillers all exhibited an obvious decline in the separation performances (Figure 44). These results were attributed to the increased membrane hydrophobicity upon integration of the AC fillers, which successfully reduced the amount of water uptake by the AC-based composite membranes. Therefore, the study clearly validates the effectiveness of ACs in tuning the hydrophobicity of the composite membranes for a stronger capacity to handle a humid biogas feed. 4.1.5.2. Carbon Molecular Sieves. CMSs constitute another class of carbon particle with a high capacity for CO2/CH4 separation given the following reasons.596−598 First, the essentially carbon-structured CMSs have attributes that are similar to MOPs and ACs. CMSs are, therefore, expected to display stronger chemical affinity and promote better interfacial adhesion with glassy polymers,429 which can potentially circumvent the need for complex strategies to mitigate the nonideal morphologies (see section 3.4.2). Second, CMSs possess micropores with constricted pore entrances that can effectively induce molecular sieving effect for CO2/CH4 separation (see section 2.1.5.2). Third, the CO 2 /CH 4 separation performance of the CMS-based composite membranes can be engineered through a facile tuning of the pyrolysis conditions.429 As such, the prospective polymer−filler compatibility, as well as tunable gas transport properties make CMSs a compelling filler for CO2/CH4 separation membranes.

due to a higher CO2 diffusivity that was induced by a narrow micropore distribution of the ACs. Following this study, efforts to enhance compatibility between the AC filler and a silicone rubber polymer matrix were carried out through the addition of plasticizers such as PEG, which generated a synergistic improvement in both the CO2 permeability and CO2/N2 selectivity.143 These optimistic results inspire the use of ACbased composite membranes for the separation of biogas. Anson et al.594 first conducted an investigation on CO2/CH4 separation using AC-based membranes. Two different commercial ACs, one with predominantly microporous domains while the other contained both micro- and mesoporous domains, were utilized in an acrylonitrilebutadiene-styrene (ABS) copolymer matrix. It was found that a partial compatibility between the AC filler and the chains of the butadiene-styrene copolymer resulted in a favorable polymer−filler interfacial contact, leading to improvements in both the CO2 permeability and CO2/CH4 selectivity. Recent studies on AC-based membranes are driven by the need to handle the presence of water vapor in the biogas. Humidified biogas from organic wastes and landfills can contain up to ∼6 vol % of water vapor, making membranebased biogas separation more challenging as compared to the natural gas (see Table 1). To meet this demand, ACs can offer the advantage of intrinsic hydrophobicity that is generally lacking in the other fillers discussed thus far. For this reason, Kanehashi et al.578,595 prepared composite membranes with a series of commercial adsorbents. This included using ACs as the filler and Matrimid as the polymer matrix. Despite a poorer separation performance of the dry CO2/CH4 mixed gas, the BE


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Figure 46. FESEM images of (a) a low and (b) a high-magnification cross-sectional morphology of the MCM-41/PSf membrane at 40 wt % loading, revealing the well-dispersed MCM-41 fillers within the polymer matrix. Reprinted with permission from ref 603. Copyright 2008 Elsevier.

filler and polymer matrix, leading to an improvement in the polymer−filler interaction. Therefore, the integration of mesoporous materials can positively enhance the CO2/CH4 separation performances of the polymeric membranes.362,602,603 To this end, mesoporous materials, including MCM-41,362 MCM-48,602 SBA-15,604 and COK-12,380 have already been utilized for this purpose. For instance, Reid et al.362 integrated mesoporous MCM-41, which had a large pore size of 40 Å, into a PSf polymer matrix for CO2/CH4 separation. A 30 wt % loading of MCM-41 provided a 3-fold enhancement in the CO2 permeability as compared to the unfilled membrane. The CO2/CH4 selectivity, however, remained unaffected despite the effective penetration of the PSf polymer chains into the pores of MCM-41.362 This was attributed to the much larger pore size of MCM-41 and the lack of strong interaction between MCM-41 and the CO2 molecules. As such, further research efforts have been subsequently focused on chemical modifications of the mesopores to improve the CO2 interaction (see section 2.1.6) and create a greater enhancement in the CO2/CH4 solubility selectivity.478 In this regard, an amine-modified mesoporous Dallas Amorphous Material-1 (DAM-1) filler loaded at 20 wt % into a Matrimid matrix demonstrated effective CO2 transport. A remarkable improvement in ideal CO2/CH4 selectivity beyond 200 was observed when using gas feeds saturated with moisture.605 Apart from this, a SPEEK matrix with 30 wt % loading of sulfonated MCM-41 showed a 20% increase in the CO2 solubility coefficient, along with ∼60 and ∼30% enhancements in the ideal and mixed CO2/CH4 selectivities, respectively, compared with an unmodified MCM41 counterpart.606 Intrinsically, mesoporous silicas like MCM-41 have poor affinity for organic polymers owing to their inorganic nature. This potentially results in a reduced polymer−filler interfacial adhesion, leading to poor mechanical properties of the composite membranes. Enabling penetration of the mesopores by polymer chains is one of the solutions known to impart greater mechanical strength to the membranes through more intimate contact.362 In addition, other solutions have been investigated to address these issues. Particularly, reducing the particle size of the mesoporous materials is a strategic approach to introduce a larger polymer−filler interfacial area to allow higher filler loading without making the composite membranes overly brittle. A study by Kim and Marand603 demonstrated this by incorporating nanosized MCM-41 of ∼80 nm to maximize the filler loading in a PSf matrix, achieving up to 40 wt % in the composite membrane. This freestanding membrane was also found to be mechanically robust but

Corroborating this, Vu et al.599,601 integrated CMSs into commercial Ultem and Matrimid polymer matrices to prepare CMS-based membranes for CO2/CH4 separation. In their studies, synergistic enhancements by as much as ∼209 and ∼46% in the CO2 permeability and CO2/CH4 selectivity were, respectively, observed in CMS-based membranes of various filler loadings. Nonetheless, the separation performances of these composite membranes continued to remain well below the upper bound limit (Figure 45a). Although the pure CMSs membrane obtained by pyrolysis of the Matrimid precursor films exhibited a permselectivity that surpassed the upper bound limit (Figure 45a), the membrane fabrication protocol was much more demanding than that of CMS-based composite membranes, owing to their inherent rigidity, brittleness, and ease of defect formation.599 Moreover, it was also capitalintensive to use CMSs as an entire continuous phase for membranes. For these reasons, pure CMSs membranes, in contrast to CMS-based composite membranes, are regarded to be of less industrial relevance. In this regard, addressing the moderate CO2/CH4 separation performances of CMS-based membranes then becomes an important research direction for strengthening the competitive position of the membranes for biogas separation. Recently, Nasir et al.600 developed a CMSbased membrane using 30 wt % loading of CMSs as the filler, PES as the polymer matrix, and DEA (5 to 15 wt %) as the third component. The addition of DEA not only promoted a better polymer−filler adhesion and improved the membrane morphology but also served as a reactive carrier to increase CO2 affinity and facilitate CO2 transport through the composite membrane. As a result, the CO2 permeability and CO2/CH4 selectivity were enhanced by a remarkable 324 and 852%, respectively, pushing the performance of such CMSbased membrane beyond the 2008 upper bound limit. Conversely, for its counterpart without DEA addition, a paler enhancement of 78 and 106%, respectively, was observed instead.600 This aptly illustrates that with careful designs and effective strategies to strengthen the CO2 affinity, separation performances can be greatly advanced to produce CMS-based membranes with attractive CO2/CH4 separation performances. 4.1.6. Mesoporous Materials. Similar to zeolites and CMSs, mesoporous materials have a long-standing history with gas separation membranes because of their well-established synthesis protocols, compositional, and morphological versatilities, as well as their well-ordered pores and tunable pore dimensions.357 In addition, their distinctively large pores, which are in the range from 2 to 50 nm, can allow smaller polymer chains of typical 1 nm in cross-sectional dimension to penetrate. This results in a more intimate contact between the BF


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Figure 47. (a) Normalized permeance of various gases using bilayer graphene membrane of different nanopore sizes before and after etching. (b) Ideal membrane selectivity of the bilayer graphene membrane with 4.9 Å nanopores. Reprinted with permission from ref 399. Copyright 2012 Nature Publishing Group. (c) TEM image of a 15 nm FIB-drilled nanopore on a freestanding bilayer graphene. (d) Gas permselectivity (defined as the permeance of H2, He, CH4, N2, CO2, or SF6 normalized by N2 permeance) of different pore diameters (7.6 to 1000 nm) with respect to molecular weight and (e) H2/CO2 separation factors against different pore diameters of the bilayer graphene membranes. Reprinted with permission from ref 389. Copyright 2014 American Association for the Advancement of Science. Plots of (f) permeance and (g) permselectivity of various gases across a few-layer GO membrane in different relative humidity. Reprinted with permission from ref 401. Copyright 2013 American Association for the Advancement of Science. (h) Single-gas permeance of various gas across a ∼18 nm GO membrane at 20 °C. Reprinted with permission from ref 607. Copyright 2013 American Association for the Advancement of Science.

various gases, which included H2, CO2, Ar, N2, CH4, and SF6, was improved with the introduction of these micropores. As the size of the micropores widened from 3.4 to 4.9 Å, the permeability of the gases increased further (Figure 47a), suggesting that the bilayer graphene membrane excluded gas via a molecular size exclusion principle (i.e., molecular sieving). The bilayer graphene membrane with a 3.4 Å pore size exhibited ultrahigh ideal CO2/N2 and CO2/CH4 selectivities of 7.0 and 9.0 × 103, respectively. With larger-sized micropores of 4.9 Å, the selectivities decreased sharply to 3.6 and 1.7, respectively (Figure 47b). A single-layer graphene with 5 Å micropores gave largely comparable selectivities at 3.9 and 4.8, respectively.399 These results provided clear evidence to support the conclusion that gas flow through the micropores was not governed by a classical effusion-based transport but rather an effusion through ångstrom-sized pores. In another study using a similar membrane design, Celebi et al. prepared a porous bilayer CVD graphene membrane with a focus-ion beam (FIB) technique to drill pores that were in the range from 7.6 to 1000 nm (Figure 47c). The gas permeance was found to exhibit no minimum Knudsen number while the single gas permselectivity presented an inverse square root proportionality with respect to the mass of the gas molecules (Figure 47d). For H2/CO2 mixed-gas separation, the selectivity displayed a pore size dependence with the membrane, which has the smallest pore size of 7.6 nm, showing the highest selectivity at 3.4 (Figure 47e).389 These findings suggested a unique gas transport phenomenon where no molecule-wall interaction occurs as a result of the near monatomic thickness of the graphene pores, and an effusive behavior, which gradually became collective when the pore sizes of the bilayer graphene membranes increased. Nevertheless, such bilayers graphene membranes are extremely difficult to fabricate and have low scalability, given the technical challenge of synthesizing large-area defect-free graphene.34

suffered from microscale voids at the polymer−filler interface due to the agglomeration of the smaller-sized MCM-41 particles. To resolve this, MCM-41 was functionalized with trimethylchlorosilane (TMCS) to render its surface more hydrophobic so as to effectively suppress the Si−Si interactions while promoting the Si-polymer interactions. Consequently, the TMCS-modified MCM-41 particles exhibited an excellent dispersion in the polymer matrix, giving a more intact membrane cross-sectional morphology (Figure 46) and a higher CO2/CH4 selectivity than that predicted by the Knudsen model.603 4.2. Two-Dimensional Materials

4.2.1. Graphene-Family Materials. GFM-based membranes for gas separation can be classified into two main categories: GFM-based (1) stacked and (2) composite membranes. GFM-based stacked membranes are asymmetric membranes comprising selective layers made up of freestanding GFMs sheets that are assembled into lamellar structures on top of the polymeric substrates. They can be subdivided according to the number of layers in the GFMs laminates (i.e., few-layer or multilayer). While each of them has their own merits and limitations, GFM-based stacked membranes, especially the few-layer ones, appear to possess extraordinary gas separation performances owing to their ultimately thin membrane thickness.389 On the other hand, GFM-based composite membranes offer the advantage of a more facile membrane fabrication process with greater versatility in terms of membrane configurations and scale-up potential. Koenig et al. first utilized a chemical vapor deposition (CVD) method to synthesize a bilayer graphene for use as a size-selective gas separation membrane. To instill the otherwise impermeable graphene membrane with a capacity to carry out selective molecular sieving, ultraviolet-induced etching of the graphene to create subnanometer pores on the membrane surface was employed. They revealed that the permeability of BG


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exfoliated, and the interlayer spacing served as well-defined nanochannels to facilitate CO2 transport through the PEBA polymer matrix (Figure 48a). On this account, a membrane

To circumvent this challenge, the water solubility of GO can be exploited to easily process the 2D nanomaterials into GO membranes using spin-casting or vacuum filtration techniques. Kim et al. proposed that by carefully controlling the stacking process via a spin-casting method, closely packed and highly interlocked GO stacked layers of ∼3 nm can be obtained on top of a support membrane. This gave rise to a higher membrane selectivity but at the expense of the gas permeability. Still, the CO2 intrinsic permeability can reach a value of ∼8500 barrer with CO2/N2 and CO2/CH4 selectivity reaching ∼20 and ∼10, respectively.401 More importantly, due to the hydrophilicity of the GO sheets, the as-prepared GO membranes exhibited reduced gas permeance (from ∼25 to ∼2 GPU) when subjected to humidity as high as 85%. This was attributed to the condensed water molecules (trapped between the interlayer spacing of the GO sheets), which hindered the transport of noncondensable small gas molecules such as H2, O2, N2, CH4, and He. Interestingly, the condensed water molecules did not seem to affect the CO2 transport as revealed by the unaffected CO2 permeance (∼100 GPU) with increasing relative humidity (Figure 47f). In fact, the GO membrane demonstrated an even higher CO2/CH4 selectivity of ∼30 at relative humidity of 85% as compared to the dry state (Figure 47g).401 These results were explained by the CO2-philic permeation being enhanced by the presence of condensed water in the interlayer spacing of the GO membrane. In a separate study, Li et al. prepared GO stacked membranes via a vacuum filtration method. GO membranes with thickness as low as 1.8 nm were prepared over an anodic aluminum oxide (AAO) membrane by optimizing the concentration of the GO solution. The study, however, revealed a contradictory result where the permeance of CO2 was lower than that of CH4 (Figure 47h).607 A possible explanation pointed toward a stronger CO2 adsorption by the oxygen-containing groups in the structural defects of the GO membranes. Although this explanation is somehow substantiated by Shen et al., who used a vacuum-spinning method to coat GO and PEI in a LbL manner, it remains controversial to date.608 In contrast with GFM-based stacked membranes, GFMbased composite membranes are comparatively easier to fabricate and scale-up due to the following reasons. First, the fabrication of conventional polymeric membranes are highly developed and industrially well-established.609 They can also be produced in hollow fiber configurations, which can offer higher packing density and are especially useful for gas separation applications (see section 3.3.2). Second, the incorporation of GFMs can be easily realized by exploiting the existing membrane fabrication technology with only moderate modifications made to the production line.34 Third, the GFM-based composite membranes are likely to be more cost-competitive as compared to the GFM-based stacked membranes because of their simpler fabrication process, commercial viability, and membrane stability.34,391 In this regard, there are more studies focused on the design of GFMbased composite membranes for gas separation. For example, Shen et al.610 prepared a GFM-based composite membrane with GO sheets in the matrix of PEBA polymer. The oxygen-containing functional groups on the GO sheets participated in hydrogen bonding with the PEBA polymer chains and gave rise to a favorable interfacial interaction, which promoted a uniform GO dispersion in the composite membrane. The integrated GO sheets were not fully

Figure 48. Schematic illustrations showing (a) the interfacial interactions between the GO filler and PEBA polymer matrix with the interlayer spacing between the sheets acting as well-defined channels for CO2 transport. Reprinted with permission from ref 610. Copyright 2015 Wiley-VCH. (b) Tortuous pathway through the polymer matrix with facilitated CO2 transport realized by the functional groups grafted on the GO sheets. Reprinted from ref 612. Copyright 2015 American Chemical Society.

containing 0.1 wt % GO exhibited a CO2 permeability of ∼100 barrer and a CO2/CH4 selectivity of 24.7,610 resulting in an enhancement of more than 100% in permeability (without compromising selectivity) as compared to the unfilled PEBA membrane. Li et al.611 also utilized GO sheets together with MWCNTs as fillers in a Matrimid matrix to create synergistic benefits for CO2 separation. Results showed that the inner channels of MWCNTs created pathways to promote CO2 transport, whereas the oxygen-containing functional groups of the GO sheets served as favorable CO2 adsorption sites, which rendered the GO sheets as effective selective barriers. Furthermore, by amalgamating two nanomaterials, the GO sheets enveloped the MWCNTs while the MWCNTs prevented the GO sheets from restacking by intercalating between them. This synergistic effect generated a more homogeneous dispersion of both fillers at 5 wt % loading each. As a result, the CO2 permeability and ideal CO2/CH4 selectivity exhibited enhancements of more than 300 and 100%, respectively, with similar enhancements observed for mixed-gas evaluations.611 The mechanical properties of the composite membranes also benefitted from this synergistic effect. With 5 wt % loading of GO sheets and MWCNTs each, the composite membranes produced a Young’s modulus of 1826 MPa, which was more than 100% increment from the unfilled Matrimid membrane and at least 5% higher than when either GO sheets or MWCNTs were solely incorporated. This increment suggested that the synergistic combination of both GO sheets and MWCNTs induced a stronger interfacial interaction with the polymer matrix, and thus significantly impacted the mechanical properties of the membranes.611 In addition, by exploiting the tunable physicochemical properties to carry out chemical modifications on the GFMs, BH


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Figure 49. (a) Mechanism of CO2/CH4 separation in a layered silicate-based composite membrane. Adapted with permission from ref 435. Copyright 2013 Elsevier. (b) Schematic illustration of a segment of the aligned montmorillonite layer, showing the interlayer spacing as an effective channel for CO2 transport. A comparison of the (c) mixed-gas permeance and (d) selectivity of the vertically and randomly aligned montmorillonite membranes. Reprinted with permission from ref 628. Copyright 2016 Wiley-VCH. (e) Different possible configurations of layered silicates in composite membranes. Adapted with permission from ref 478. Copyright 2013 Royal Society of Chemistry.

in the SPEEK polymer matrix became longer and more tortuous with GO incorporation. Moreover, in the humidified state, the water that was trapped within the SPEEK matrix facilitated a lower transport resistance for CO2 but not for CH4 given its limited solubility in water. The CO2 facilitated transport was also intensified by the presence of the amine reactive carriers grafted on the GO sheets. As such, combining these effects increased the solubility and diffusivity of CO2, which thus led to synergistic enhancements in the CO2 permeability and CO2/CH4 selectivity of the humidified membrane.613 Similarly, Li et al.612 chemically grafted GO sheets with PEI and PEG prior to their integration within a Pebax matrix. The design capitalized on the grafted ethylene oxide and amine functional groups as reactive sites to enhance the CO2 affinity and transport, while simultaneously increasing

the CO 2 /CH 4 separation performance of GFM-based composite membranes can be further engineered and optimized. For instance, Xin et al.613 chemically grafted GO sheets with dopamine and cysteine before incorporating the functionalized GO sheets into a SPEEK polymer matrix. Due to the effect of the amino acids functional groups on the GO sheets, the solubility of CO2 molecules increased, and this enhanced the CO2 permeability close to 45%, reaching a value of 22.3 barrer. In particular, when the membrane was humidified to a water content of ∼27%, the CO2 permeability achieved a high value of ∼1220 barrer without compromising the membrane selectivity. The CO2/CH4 selectivity in fact increased from ∼50 to ∼80 when the membrane state changed from dried to humidified.613 In agreement with previous discussion, the diffusive pathway of the larger CH4 molecules BI


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the cations that are already present in the gallery.624,625 Accordingly, these species exhibit great capacities to intercalate into the layered silicates, causing expansion to the interlayer spacing. Nonetheless, it is very challenging to utilize a generalized approach to induce swelling of the layered silicates given a wide variation in their structures. For instance, the swelling of AMH-3 interlayer spacing is hampered by the presence of strong ionic bonds between Na+ and Sr2+ cations and the silicate layers. These ionic bonds are difficult to overcome by hydrogen bonding, not to mention the risk of interlayer condensation (which forms strong covalent bonds between layers) if cation-exchange reaction is employed instead. Therefore, to introduce swelling in AMH-3, Choi et al.626 initiated a novel approach involving sequential intercalation of dodecylamine after an ion-exchange process using an amino acid, DL-histidine. The procedure was later refined to preserve the overall pore structure of AMH-3 to a greater extent during the swelling process.627 This allowed the 2D layered silicate to put to advantage its 3.4 Å pore aperture for CO2/CH4 membrane-based separation. As a result, the composite membrane using a 6 wt % loading of swollen AMH3 in a CA polymer matrix exhibited a 54% enhancement in the CO2 permeability without compromising the CO2/CH4 selectivity as compared to the unfilled CA membrane.435 Third, the alignment of the lamellar structure of the layered silicates can critically impact the CO2 transport properties of the composite membranes. To demonstrate this, Qiao et al.628 exploited the uncoiled chain stretching orientation of polyvinylamineacid and the coupling effect of APTES to bond and align montmorillonite (MMT) layers vertically onto a porous PSf substrate (Figure 49b). In a 1:9 CO2/CH4 mixedgas evaluation, the separation performances of the vertically aligned MMT membrane showed an 83% increment in the CO2 permeability at a comparable CO2/CH4 selectivity, in contrast to its randomly aligned counterpart (Figure 49c,d). The largely similar selectivity indicated a robust alignment approach, which resulted in the formation of a defect-free selective layer. The higher permeability, on the other hand, suggested a high-speed CO2 transport owing to the more direct channels created by the vertical interlayer spacing of the MMT. In comparison, randomly aligned membranes offered a longer diffusive pathway due to the random orientation of the MMT and a higher mass transport resistance given the presence of a continuous phase that was made up of polyvinylamineacid. Thus, the CO2 permeability of these two membranes were distinctively different. Taken together, we believe that the success of layered silicate-based composite membranes is closely associated with the interfacial compatibility, uniform dispersion, proper alignment, and sufficient exfoliation of the nanoporous layered silicates. Usually, defective layered silicate-based membranes with unoptimized CO2/CH4 separation performances are due to either poor polymer−filler matching or inappropriate membrane fabrications, which result in agglomerated, unaligned, or unexfoliated layered silicates within the composite membranes (Figure 49e). Hence, future research in silicatebased composite membranes should be focused on solutionbased intervention strategies to avoid such pitfalls.

the tortuosity of the CH4 diffusive pathway using impermeable GO sheets of high aspect ratio (Figure 48b). For these reasons, the composite membrane exhibited a CO2 permeability of 145 and 1330 barrer, whereas the CO2/CH4 selectivity improved to 24.0 and 45.0 under a dried and humidified state, respectively.612 Overall, these studies exemplify the strategic use of GO nanomaterials as versatile fillers to introduce specific functional groups for carrying out CO2 facilitated (carrier mediated) transport within the polymer matrix. 4.2.2. Layered (Lamellar) Silicates. The use of layered silicates in polymeric membranes has been driven mainly by the vast accumulated knowledge on intercalation chemistry, as well as the industrial interest and academic value they create in terms of their abilities to improve the mechanical, thermal, and barrier properties of the polymeric materials.614,615 Compared to the conventional fillers (such as calcium carbonate, glass fibers, and mica), layered silicates or clays are capable of dispersing effectively within the polymer matrices, resulting in less defective and brittle composite membranes. 428,429 Furthermore, due to their 2D structures, layered silicates can exhibit an aspect ratio as high as 1000 when fully exfoliated into individual sheets. Thus, low filler loadings will usually suffice for most composite membranes. Layered silicates in composite membranes are traditionally employed in the food packaging industry where they are utilized to increase the oxygen barrier properties of the packaging films.616,617 Novel applications like gas separation are also starting to gain increasing attention, owing to the costeffectiveness, abundant availability, and greener manufacturing process of layered silicates.618 More importantly, nanoporous layered silicates demonstrate a similar barrier effect like 2D GFMs fillers, which is capable of altering the gas transport properties of membranes by creating a tortuous diffusive pathway for the larger CH4 molecules, while enabling a more direct pathway for the smaller CO2 molecules (Figure 49a). Despite this, there are currently limited number of studies on layered silicate-based composite membranes for CO2/CH4 separation in the literature. Instead, more extensive studies on H2-based separation using JDF-L1436 and MCM-22619,620 as fillers have been reported. Therefore, we believe that there is still room for expansion in this area of research. Strategically, a key focus for layered silicate-based composite membranes is to resolve existing issues with the use of layered silicates. First, the layered silicates are mostly hydrophilic, making them less compatible with polymeric matrices that are commonly used for membrane preparations. To alleviate this, surface modifications such as grafting of silane coupling agents have yielded good results as previously described in Section 3.4.2.1. Furthermore, the silane agents with positively charged amine functional groups can imbue a strong repulsive force between the layered silicates and enable uniform dispersion of the silicate fillers in the polymer matrix.621 The amine groups can also act as reactive carriers to enhance CO2 affinity of the composite membranes. Second, the layered silicates need to be reasonably swelled in order for the interlayer spacing to expand and act as effective molecular sieves for CO2/CH4 separation. This is generally observed when layered silicates undergo intercalation with suitable amines or alkyl ammonium salts. The amine molecules can diffuse into the interlayer spacing by forming hydrogen bonds with the surface silanol groups in the gallery of the silicate layers.622,623 On the other hand, the diffusion of alkyl ammonium cations is driven by an ion-exchange reaction with

4.3. One-Dimensional Materials: Carbon Nanotubes

Analogous to the aforementioned 2D material-based membranes, CNT-based membranes can also be prepared by integrating CNTs either within the selective layer or the BJ


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Figure 50. Schematic illustration of the fabrication of (a) a vertically aligned MWCNT membrane in polystyrene matrix. Reprinted with permission from ref 629. Copyright 2012 Elsevier. (b) Vertically aligned DWCNT membrane with step 1−2, substrate preparation; step 3, DWCNTs growth; step 4−5, silicon nitride encapsulation and membrane area definition; step 6−7, ion milling and etching to open up the CNTs. (c) Gas selectivity of the DWCNT membrane defined as permeability relative to He. The triangle and circle symbols represent data of DWCNT (∼2 nm) and MWCNT (∼7 nm) membrane, respectively. Open symbols denote data for nonhydrocarbon gases, while solid symbols are for hydrocarbon gases. The solid and dotted lines are power-law fit with exponential values of −0.49 and −0.37, respectively. The inset shows the full mass range of the nonhydrocarbon gases. Reprinted with permission from ref 443. Copyright 2006 The American Association for the Advancement of Science. FESEM images of the vertically aligned MWCNT array (d) before and (e) after polystyrene encapsulation. Reprinted with permission from ref 453. Copyright 2004 The American Association for the Advancement of Science.

the inner channels of

Harnessing Filler Materials for Enhancing Biogas Separation

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