Rigid and microporous polymers for gas separation membranes

Progress in Polymer Science 43 (2015) 1–32

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Progress in Polymer Science journal homepage: www.elsevier.com/locate/ppolysci

Rigid and microporous polymers for gas separation membranes Seungju Kim, Young Moo Lee ∗ Department of Energy Engineering, College of Engineering, Hanyang University, Seoul 133-791, Republic of Korea

a r t i c l e

i n f o

Article history: Received 11 October 2013 Received in revised form 21 October 2014 Accepted 22 October 2014 Available online 1 November 2014 Keywords: Microporous polymers High performance polymer membranes CO2 separation Rigid polymers

a b s t r a c t Microporous polymers are a class of microporous materials with high free volume elements and large surface areas. Microporous polymers have received much attention for various applications in gas separation, gas storage, and for clean energy resources due to their easy processability for mass production, as well as microporosity for high performance. This review describes recent research trends of microporous polymers in various energy related applications, especially for gas separations and gas storages. The new classes of microporous polymers, so-called thermally rearranged (TR) polymers and polymers of intrinsic microporosity (PIMs), have been developed by enhancing polymer rigidity to improve microporosity with sufficient free volume sizes. Their rigidity improves separation performance and efficiency with extraordinary gas permeability. Moreover, their solubility in organic solvents allows them to have potential use in large-scale industrial applications. © 2014 Elsevier Ltd. All rights reserved.

Contents 1. 2. 3.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Membrane gas separation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Rigid and microporous polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Microporous polymers as energy saving materials. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1.1. Microporous polymers for gas storage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1.2. Microporous polymers for gas separation membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Thermally rearranged (TR) polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1. TR polymers derived from functionalized polyimides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.2. TR polymers derived from polyamides with hydroxyl groups . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.3. Polyimides with thermally labile units . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.4. TR co-polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.5. Physical property and characterization of TR polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. Polymers of intrinsic microporosity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.1. Ethanoanthracene (EA) and Tröger’s base (TB) based PIMs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.2. Cross-linked PIMs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.3. PIMs with substituted pendant groups . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

∗ Corresponding author. Tel.: +822 2220 0525. E-mail address: [email protected] (Y.M. Lee). http://dx.doi.org/10.1016/j.progpolymsci.2014.10.005 0079-6700/© 2014 Elsevier Ltd. All rights reserved.

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3.3.4. PIM blends and mixed matrix membranes (MMMs) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.5. Polyimides of intrinsic microporosity (PI-PIMs) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.6. TR polymers with spirobisindane (PIM-TR-PBO) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Future outlooks and conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction Microporous materials have been of great interest for many potential applications in fields such as clean energy, catalysis, and storage media, due to their extraordinarily high porosity and surface area. Microporous organic materials, such as various network polymers, have emerged in applications including gas separation, gas storage, light harvesting, energy storage, and energy conversion [1–4]. Moreover, the requirement of clean and alternative energy resources, such as hydrogen energy, fuel cells, and energy storage, has been growing with the increase of worldwide energy usage [2–18]. Following the recent definition from the International Union of Pure and Applied Chemistry (IUPAC), microporous materials involve extremely small pores with average sizes of less than 2 nm [19]. These small pores in microporous materials induce transport and storage of small molecules and ions. Hydrogen energy has been considered as a clean energy resource, whereas carbon dioxide (CO2 ) capture technology has been developed to reduce the emission of greenhouse gases to the atmosphere [20–23]. In these applications, microporous materials including microporous organic polymers (MOPs), zeolitic imidazolate frameworks (ZIFs), and metal organic frameworks (MOFs) have been developed with a high porosity and large surface area as gas storage and separation materials. Moreover, the easy processability of microporous polymers has demonstrated a great advantage for expansion into industrial applications [11,24,25]. Hydrogen resources possess a high energy density compared with hydrocarbon-based resources. Combustion residuals from hydrogen are only clean water. However, many issues in hydrogen production, delivery, and storage in an economic and safe way limit the growth of hydrogen energy uses. Energy conversion devices for hydrogen, such as fuel cell systems, have been developed for both mobile and stationary applications. However, a successful hydrogen resource should satisfy the production, delivery, and storage issues [14,25]. The lack of convenience, safety, and cost effectiveness in hydrogen storage are major challenges, and thus various materials have been investigated to maximize hydrogen uptake capacity in storage materials [22,26]. For CO2 separation, removal of CO2 as an impurity is the largest challenge in the natural gas treatment, landfill biogas recovery, and enhanced oil recovery (EOR) processes [15,27–29]. More recently, CO2 has been considered as one of the most influencing greenhouse gases for global warming. So far, microporous materials have been investigated as CO2 absorbents and separation membranes [1,11,30–32]. Carbon capture and storage (CCS) technology has been widely researched from the separation and sequestration processes, as well as CO2 conversion into valuable

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resources [33,34]. CO2 emission during power generation is a major emission source, which includes fuel combustion (post-combustion, CO2 /N2 separation) and the water gas shift (WGS) reaction in the integrated gasification combined cycle (IGCC) process (pre-combustion, H2 /CO2 separation). IGCC process has been applied to turn coals into gases for clean thermal power generation [35,36]. The requirement of CO2 separation from various emission sources as a carbon capture technology is growing [35]. The separation process works against the law of thermodynamics, as a result, separation technology involves great deal of energy consumption [37]. The requirement for separation and purification technologies includes productivity and purity, which are directly related to the process cost. Therefore, the quality of the separation technology depends on the choice of separation process, as well as the process design to achieve a high productivity and purity [37]. The traditional gas separation technologies, i.e., absorption, adsorption, and cryogenic distillation, require a great deal of energy consumption as well as a large site area to produce pure gases [38]. Membrane gas separation has played an important role with its low energy usage compared to traditional separation technologies with increasing energy cost [29,39]. Membrane gas separation processes do not require phase change or an extra thermal regeneration process, so they have a potential of energy efficiency competitive with other separation processes. Membrane gas separation processes also require relatively small footprints, which reduce the operation site requirements [28]. Various materials such as polymers, metals, ceramics, and hybrid materials have been used as membrane materials. Among membrane materials, polymers are dominant in large scale applications due to their easy processability, as well as good mechanical property. Polymers could be formed to an asymmetric structure, which can be used to fabricate high flux membrane modules for large scale applications [40]. Various polymeric membrane modules have been commercialized by several companies, but only a small number of commercial polymers such as cellulose acetate (CA), polyimides (PI), and poly(phenylene oxide) (PPO) have been used. Gas permeability and selectivity represent membrane performance. Gas permeability is a term of productivity and gas selectivity refers to product purity [41]. Most commercial gas separation membranes are produced from low permeability polymers with reasonable selectivity. Therefore, a number of membrane modules are required for high productivity [42]. The tradeoff relationship of gas separation membranes between gas permeability and selectivity is well-defined, where highly permeable materials usually show low separation properties and vice versa. The trade-off relationship between permeability and selectivity is mainly due to permeation

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efficiency since high permeability membranes usually possesses either many diffusive pathways or sorption sites, where other gases can also pass through, and as a result, such membranes offer either poor efficiency or gas selectivity [43,44]. According to a report from the National Energy Technology Laboratory in the U.S. Department of Energy (NETL, DOE), 90% CO2 capture from power plants with a 35% increase in the cost of electricity (COE) are predicted by 2030 [45–47]. In this area, CO2 capture using monoethanolamine (MEA) is the most investigated and mature technology, but the associated process cost is high. MEA would require a CO2 capture cost of $40–100/tonCO2 [47]. However, membrane CO2 capture will satisfy the DOE target by development of highly permeable membrane materials. According to Merkel et al., their membrane system, where CO2 permeance is about 1000 GPU (1 GPU = 10−6 cm3 (STP)/cm2 s cmHg = 3.35 × 10−10 mol/m2 s Pa) and CO2 /N2 selectivity is 50, could reduce the CO2 capture cost to $23/tonCO2 . They also emphasized the importance of highly permeable materials where over 2000 GPU for CO2 permeance with a CO2 /N2 selectivity over 25 would achieve a CO2 capture cost less than $20/tonCO2 [46]. Over the last 40 years, substantial studies on polymer membrane materials have been performed to improve their performances. Consequently, both permeability and selectivity have been improved compared with first generation materials [42,48–50]. Current membrane performance has a limitation, which is called the “upper bound”. From a large number of data in previous references, Robeson suggested an “upper bound” in the trade-off between gas permeability and selectivity for valuable gas pairs. After 17 years, the upper-bound was revisited with state-of-the-art data [43,44]. The development of highly permeable membranes with acceptable selectivity can significantly reduce the process costs. Product purity, related with gas selectivity, can be achieved by process design. However, productivity can only be achieved by a number of membrane modules with a large area, or the use of highly gas permeable membrane materials [46]. Therefore, the breakthrough in the development of highly permeable membrane materials is essential. Microporous polymers have been considered to overcome the upper-bound as efficient membrane materials. To date, microporous polymers have been mainly studied as storage materials or sorbents. However, their high free volume elements could be advantageous as gas separation membranes with high gas permeability [48]. Microporous polymeric membranes, such as substituted polyacetylenes and amorphous fluoropolymers, have demonstrated higher permeability than other conventional low free volume polymers. However, due to the trade-off relationship, they exhibit a low selectivity [48]. In the last decade, a new class of microporous polymers with sufficient rigidity, such as thermally rearranged (TR) polymers and polymers of intrinsic microporosity (PIMs), has been reported to have high selectivity, as well as extraordinary gas permeability compared with other membrane materials such as CA, PI, and PPO. The membrane performance of these microporous polymers also surpassed the upper-bound [10,51]. They have provided the potential for a more energy efficient

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membrane gas separation process in many industrial applications. In this review, we present the research progress on rigid and microporous polymers as gas storage materials, as well as gas separation membranes. 2. Membrane gas separation Membrane gas separation is a pressure-driven process by which the pressure difference between upstream and downstream of the membrane acts as a driving force. Gas separation membranes are usually known as non-porous materials where the pore sizes are extremely small, so the pores are not visible and a special method is required to characterize the pores. Gas transport through the polymer membrane is based on a solution–diffusion mechanism where the gas transport rate is related with the affinity between the gas molecules and membrane materials, as well as the gas diffusion rate through the membrane matrix. Permeation of gas molecules through membranes is explained by a solution–diffusion mechanism, which is described as P =D·S

(1)

where P, D, and S are permeability, diffusivity, and solubility coefficients, respectively. Gas permeability is considered as a product of diffusivity and solubility. Permeability unit is commonly represented as Barrer (1 Barrer = 10−10 cm3 (STP) cm/cm2 s cmHg = 3.35 × 10−16 mol m/m2 s Pa) to evaluate membrane material performances, whereas permeance, which is defined as thickness irrelevant permeability (P/L), is considered to estimate membrane module performances or asymmetric or hollow fiber membrane performance with a unit of GPU. Gas diffusion is correlated to the free volume in polymers and the size of penetrating gases, whereas gas solubility is related to the chemical affinity between gas molecules and the polymer matrix. Ideal selectivity is considered as the ratio of permeability coefficients. a=

Pi D S = i · i Pj Dj Sj

(2)

where ˛i/j is an ideal selectivity in a gas mixture of i and j [39,40,52]. Diffusion is a kinetic factor that is related with the spatial extents of random motion in a bulk polymer matrix. On the other hand, sorption is a thermodynamic factor that occurs when gas molecules physically adsorb onto a membrane surface and in a bulk polymer matrix with respect to the concentration difference [40]. Sorption and diffusion are also activated energy processes, so they are frequently affected by temperature. Polymer types determine the affinity between gas penetrants and polymer matrix which makes difference in solubility. Therefore, solubility and diffusivity coefficients are significantly considered when developing membrane materials. Two types of gas separation membranes, diffusion-selective and sorptionselective, are known where the membrane performance can be improved by enhancing the diffusivity and/or solubility [52]. Diffusion-selective membranes are usually glassy polymer membranes with a high glass transition

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temperature (Tg ) and high free volume. Gas diffusivity and diffusion selectivity in polymer membranes are influenced by the chain rigidity and free volume elements. Glassy polymers are usually affected by the diffusion selectivity due to the dominant gas diffusion through the free space between polymer chains, the so-called free volume or cavity [53–55]. Glassy polymers including polyimide (PI), cellulose acetate (CA), and polysulfone (PSF) are traditional gas separation membranes that have been widely used over the past three decades [29,41,42,48]. Those polymers have low free volume elements. Consequently, they demonstrated low gas permeability, but high selectivity even at high operating pressure. Moreover, glassy polymers are easily affected by plasticization, that decreases gas selectivity at high pressure and spoils the process efficiency [42]. Sorption-selective membranes are usually rubbery polymer membranes with low diffusion selectivity. Rubbery polymers, which have a low Tg below room temperature, demonstrate high chemical affinity between polymer chains and gas molecules due to the flexible chain motion in the rubbery state [56,57]. They also have advantages including remarkable permeability for condensable gas molecules such as carbon dioxide, hydrocarbons, and volatile organic solvents (VOSs), because they are easily dissolved in the polymer matrix. Poly(dimethyl siloxane) (PDMS), poly(ethylene oxide) (PEO)-based polymers, poly(amide-6-b-ethylene oxide) (Pebax® ), polyvinylamine (PVAm), and polyvinyl alcohol (PVA) are representative examples of sorption-selective membranes [58–66]. However, according to Robeson’s reports, most rubbery polymer membranes exhibited lower gas separation performances than glassy polymer membranes for small gas pairs [43,44]. Various modification methods have been researched to prevent the plasticization effect on rubbery polymer membranes, such as crosslinking and blending [67–72]. The latest studies have focused on the microporous materials with high free volume elements, which will increase gas permeability by enhancing gas diffusivity. Microporous polymers have received much attention as membrane materials, which improve the gas transport performances and have a potential for real application. 3. Rigid and microporous polymers 3.1. Microporous polymers as energy saving materials Microporous materials have been an important issue in a wide range of applications, including catalysis, separation, and storage as energy saving materials for environmental problems [11,18,19,24]. Microporous materials can reduce energy consumption when they are applied for various processes, such as gas separation or chemical reaction, as alternatives for traditional materials. That is why microporous materials are considered as energy saving materials [48,73]. Various types of inorganic or hybrid microporous materials, such as zeolites, activated carbons, silica, and MOFs, have received much attention due to their large surface area [74,75]. However, microporous organic materials have lately been considered with their potential advantages of synthetic diversity, relatively low

cost, and better processability for mass production [19]. Microporous polymers possess interconnected pores on a nanometer and/or even sub-nanometer scale with ultrahigh porosity which is normally caused by rigid structure in polymer chain or continuous network of interconnected intermolecular voids [11,18]. They also exhibited beneficial growth in a variety of industrial applications and academic interests, such as molecular storage, separation, delivery, and catalysis [11,18,19,74,76–85]. The performances in these applications, especially in gas storage and gas separation, depend on the presence of a large surface area and a high free volume element, which are comparable with inorganic microporous materials [78–80,86–89]. Various types of microporous organic materials have been studied including porous coordination polymers (PCPs) [7,90–94], hypercrosslinked polymers (HCPs) [26,80–83,85], conjugated microporous polymers (CMPs) [2–5,89,95], and covalent organic frameworks (COFs) [20,24,96–98]. Fig. 1 shows representative structures of microporous polymers. 3.1.1. Microporous polymers for gas storage Microporous polymers as an energy saving material have mainly been investigated as storage materials for hydrogen and methane, as well as carbon dioxide capture applications [18,24]. For hydrogen and methane storage application, development of high capacity storage materials is still a major challenge. For hydrogen storage materials, the U.S. DOE has targeted a 5.5 wt% in gravimetric capacity, i.e., 400 g/L of volumetric capacity at 40–60 ◦ C under a maximum of 100 atm [25]. The potential to satisfy the target is based on the microporous materials with a large surface area and appropriate pore sizes for strong affinity to gas molecules [24]. For CO2 capture application, high CO2 uptake as well as CO2 selective adsorption ability with N2 or CH4 is important [24,75]. Various microporous organic materials have been investigated for CO2 capture applications. Microporous polymers demonstrated high performance for CO2 captures application based on their large surface area with microporosity. CMPs have been investigated since the discovery of electrically conductive-conjugated polyacetylene. CMPs include the polymer structure of linear polymers, dendrimers, hyperbranched polymers, and networks with a conjugated structure such as poly(aryleneethynylene), poly(phenylene butadiynylene), poly(p-phenylene), and polysilane networks [3]. The conjugated polymer chains were demonstrated as rigid and inflexible, but having a highly porous polymer structure. Dawson et al. have investigated various synthetic CMPs for CO2 capture application under post-combustion conditions [20]. Their CMPs showed a high CO2 uptake of 3.36 mmol/g at 1 bar and 25 ◦ C, and a potential for scalability. For high-performance CO2 adsorbents, several representative properties are required: CO2 uptake capacity (mmol/g), regenerability (%), and selectivity [21]. In selective CO2 capture by microporous adsorbents, isosteric heat to increase affinity between materials and CO2 molecules is important. Various modification methods such as pore size tailoring, metal doping, introducing functional groups, have been considered to develop material performance as CO2 absorbents [24]. Metal-coordinated CMPs have been investigated by Xie

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Fig. 1. Examples of microporous polymers mainly studied for gas storage applications: (a) [Cd(4-btapa)2 (NO3 )2 ]·6H2 O·2DMF-based porous coordination polymers (PCPs) [94], (b) conjugated microporous polymers (CMPs) [89], and (c) covalent organic frameworks (COFs) [97]. Sources: [94], Copyright 2007; [89], Copyright 2008; [97], Copyright 2009. Reproduced with permission from the American Chemical Society.

et al. [1]. Cobalt or aluminum-coordinated CMPs exhibit reversible and outstanding CO2 adsorption capacities, which are comparable with MOFs. CO2 uptake exhibited 79.3 mg/g and 76.5 mg/g in Co-CMP and Al-CMP with a BET surface area of 965 m2 /g and 798 m2 /g, respectively, whereas CMP absorbed 71.0 mg/g. These coordinated CMPs also displayed excellent catalytic activities in the conversion of CO2 and propylene oxide to propylene carbonate [1]. PCPs are flexible and dynamic porous frameworks, which are constructed from transition metal ions and bridging organic ligands with a high degree of porosity [92,93]. They have been studied as an adsorbent for gases and supercritical gases. Their porosity is controllable using a different bridging ligand and metal ions. One example of PCP materials is a 3D Cu coordination polymer with 1D channel, where the polymer networks from these coordination polymers provided open lattice species and various topologies with larger pores [91]. Various approaches to increase the porosity in coordination polymers have been studied and their application for storage has been investigated [92,93]. HCPs have widely been investigated since 3D polystyrene networks were initially reported in 1990 [99]. The most well-known HCPs are polystyrene-related polymers, so-called “Davankov-type” resins. They have been prepared by post-crosslinking, as well as by direct polycondensation. Moreover, various structures have been proposed to prepare HCPs including polyphenylenes, and their application is focused on gas storage, such as hydrogen and methane [22,26,80–83,85,99]. COFs are entirely constructed with light organic elements linked by strong covalent bonds [96,97]. Their strong bonds maintain a high porosity with 2D or 3D building blocks. The building dimensions provide the property of COFs, and as high porous materials, 3D COFs extended the framework three dimensionally thereby yielding high surface areas with low density. 3.1.2. Microporous polymers for gas separation membranes Microporous polymers have widely been investigated as gas storage materials but, applications for gas separation membranes have been quite limited. Most of microporous materials have formed insoluble networks

due to their rigid and cross-linked structure [3]. Therefore, their insolubility limits the use for more interesting applications based on their microporosity because it brings out difficulty in membrane manufacturing process [100]. For gas separation membranes, both microporosity and membrane processability are important requirements. Solution-processability is especially significant to fabricate microporous materials to any kinds of membrane types for gas separation applications. Therefore, only a few examples of soluble microporous polymers have been investigated [10,100–102]. Appropriate pore sizes and size distribution with a high porosity are desired for efficient separation performances with high gas permeability [10,103]. Therefore, membrane gas separation application using microporous organic materials has recently been reported. Microporous polymers as gas separation membranes, such as poly(trimethylsilyl-1-propyne) (PTMSP), have been reported to improve membrane performance in terms of gas permeability. However, PTMSP does not show efficient separation properties for small gas molecules [104,105] although their gas permeability is the highest among polymeric membranes. Their separation performance is not sufficient for efficient gas separation, because the polymer main chain structure is not rigid enough to retain sufficient separation properties. AF polymers are fluoropolymers with large free volume elements. They have been investigated to reduce crystallinity and to increase fractional free volume. Commercially available Teflon AF 2400 and AF 1600 are the most developed AF polymers with high gas permeability. However, their relatively flexible polymer backbone structure prevents adequate gas separation performances [106]. Polymer stiffness or rigidity in the main chain can improve both the gas permeability and selectivity. A stiff backbone, which enhances diffusional selectivity, could form the molecular windows of well-tuned sizes, while also disrupting inter-chain packing to increase permeability [107,108]. Recent progress in microporous polymers development as gas separation membranes has been investigated by enhancing the entire polymer structure rigidity to improve separation performance. A rigid structure improves the separation property and durability of membranes as gas separation and storage materials [108]. Thermally rearranged (TR) polymers and polymers of intrinsic

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Fig. 2. Thermal rearrangement mechanism, (a) TR-␣ polymer and (b) TR-␤ polymer.

microporosity (PIMs) are examples of microporous polymers with high permeability and selectivity for gas pairs in useful applications [103,109–111]. They have demonstrated an extraordinary gas permeability and selectivity with large surface area, free volume elements, and welltuned pore sizes and size distributions. The use of PIMs was known for gas storage applications at first [112–115]. However, their gas permeation and separation performance as membrane materials are advanced compared with other conventional membrane materials. Therefore, recent studies on PIMs as membrane materials have received much attention due to outstanding gas permeation and separation performances [116–118]. 3.2. Thermally rearranged (TR) polymers Gas molecules permeate through the so-called cavities or free volume elements of the polymer matrix. Therefore, microporous membranes with high free volume elements can achieve fast gas transport. The size difference between the gas molecules and the membrane cavities enables a membrane to separate gas molecules. However, because the cavity sizes in polymeric membranes usually exhibit a broad distribution, separation performance is reduced. Recently, rigid polymer membranes derived from functionalized polyimides by a post-thermal conversion process have been developed, so-called thermally rearranged (TR) polymers [51]. TR polymers are

examples of novel membrane materials with high free volume elements and narrow cavity size distributions based on rigid and microporous structure. TR polymer membranes demonstrate extraordinary gas permeability with reasonable selectivity. Their potential applications in gas separation have been widely studied. Moreover, due to their microporous properties, other applications, such as pervaporation, are also interesting [119,120]. Pervaporation is a membrane based process where the liquid feed mixture is contacted at one side of the membrane while the permeate side is applied for vacuum or sweep for efficient separation. Pervaporation process has been used to separate various liquid mixtures especially for azeotropic mixtures [119]. One of the great benefits of TR polymers is the ability to control the cavity sizes by designing various polymer structures and thermal reaction mechanisms for specific gas separation applications, including CO2 capture from flue gas [51]. Thermal rearrangement is a thermal reaction where precursor polymer is converted into other rigid structures in the solid state, as shown in Fig. 2. The precursor polymer is basically functionalized polyimides and their TR polymers are chemically and thermally stable aromatic polymers with heterocyclic rings including polybenzoxazoles (PBO), polybenzothiazoles (PBZ), polypyrrolone (PPL), and polybenzimidazoles (PBI) depending on the functional groups at the ortho position. The rigid chemical structure of TR polymers demonstrates outstanding

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7

Fig. 3. Commercial and synthetic monomers for synthesis of TR polymers and their copolymers: (a) dianhydrides, (b) acid chlorides, (c) hydroxyl diamines, (d) diamines and tetramine, and (e) anhydride chloride.

chemical and thermal resistances, which are important attributes for harsh and aggressive operation environments [51,121–134]. Thermal rearrangement process results in the evolution of interconnected micropores with a narrow size distribution, which induces rapid gas transport with size sieving separation. As a result of thermal rearrangement, the gas permeability of TR polymer membranes are improved at least two orders of magnitude over their precursor polymers, as well as most typical glassy polymers. The gas permeability of TR polymer membranes are lower than that of PTMSP, but the gas selectivity for CO2 separation is two to three fold higher than that of PTMSP, which is the most efficient gas separation possible from TR polymer membranes. Thermal rearrangement process has been extended into thermal conversion from hydroxyl group-containing polyamide precursors (HPAs) to TR polymers with a PBO structure [135,136]. One of the other strategies in the TR process is an introduction of thermally labile molecules on a cross-linkable polyimide to prepare highly permeable polyimide membranes by thermal decomposition of labile units in a solid state [137–140]. They have been investigated to control the gas separation performance for particular gas applications by tuning the cavity sizes and size distributions resulted from thermal reactions. Those membranes exhibit different gas permeability and selectivity properties that can be exploited in gas separation applications, such as H2 /CO2 , CO2 /N2 , CO2 /CH4 , and O2 /N2 separation. 3.2.1. TR polymers derived from functionalized polyimides The precursors of TR polymers are functionalized polyimide and polyamide. In the case of the TR polymer made

from functionalized polyimide, called TR-␣ polymer, the precursors were prepared by conventional polycondensation reaction of dianhydride and diamine with functional groups at the ortho position, as described in Fig. 3(a) and (c), in common polar aprotic solvents, such as N-methyl2-pyrrolidone (NMP). Different monomer structure leads to different precursor polyimides as well as the resulting TR polymers after thermal rearrangement. Precursor polyimides were obtained by various imidization methods including thermal, chemical, ester-acid, and azeotropic imidization based on a dehydration reaction [51,121,129]. The final TR polymer membranes were formed after thermal rearrangement of the precursors, prepared by a solution-process method above 350 ◦ C under inert atmospheric conditions. Thermal rearrangement is basically a thermal conversion reaction in the solid state where the precursor polyimide is converted to TR polymer as shown in Fig. 4. Therefore, thermal rearrangement conditions, such as temperature and time, for complete conversion to TR polymers are varied depending on the chemical structure and imidization method of the precursor polyimide. Thermal rearrangement or conversion temperature (TTR ) is usually determined by thermogravimetric analysis (TGA) comparing experimental weight loss with theoretical weight loss, and then confirmed by Fourier transform infrared (FT-IR) spectroscopy. Generally, precursor polymers with a high Tg demonstrated high TTR . Crosslinking occurs at high temperature during the thermal conversion where crosslinking can be determined by gel fraction measurement, which will be discussed in the following section [130]. Moreover, thermal rearrangement process is not in a linear relationship between thermal conversion ratio and TTR , where thermal rearrangement is accelerated at high temperature as schematically shown in Fig. 5.

8

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

Fig. 4. Thermal rearrangement mechanism from hydroxyl polyimide to polybenzoxazole (TR-␣-PBO) [122].

The most widely studied TR polymer is TR-PBO obtained by thermal rearrangement of hydroxyl polyimides (HPI). TR polymer membranes demonstrated a distinguished gas transport behavior depending on the polymer backbone structure, as well as the preparation method [121,129]. Gas transport properties of TR polymer membranes are also dependent on the degree of thermal rearrangement, which can be controlled by thermal reaction time and conditions [122]. They are also dependent on the chemical structure of precursor polymers, which are varied via the choice of diamines and dianhydrides [124–127]. Generally, monomers with a rigid backbone produce TR polymer membranes with high permeability. The gas transport of TR polymer membranes exhibits enhanced permeability, especially for small gas molecules, such as H2 and CO2 . The CO2 permeability of TR-PBO membranes reaches up to 5000 Barrers from a few hundred Barrers with a CO2 /N2 selectivity of 35 and CO2 /CH4 selectivity of 55, depending on the imidization method of precursors [128]. The imidization method of precursors plays important roles in the physical properties of TR-PBOs, because it

Fig. 5. Schematic of typical thermal rearrangement or conversion ratio as a function of treatment temperature.

affects the chemical structures of precursor polyimides, as well as the free volume elements of the resulting TR polymers. Han et al. investigated the gas transport properties of TR polymer membranes according to the imidization method of precursor polymers. Precursor polyimides can be prepared by thermal, chemical, and azeotropic imidization methods. Polyimides, prepared from thermal and azeotropic imidization, introduced hydroxyl-containing polyimides (HPI), whereas chemical imidization brought acetate-containing polyimide (AcPI). The bulky side chain of the acetate groups in AcPI increased free volume elements in the polymer matrix and resulting TR-PBO from AcPI, which was formed from the deformation of bulky acetate groups and decarboxylation, demonstrated larger free volume elements than TR-PBO from HPI. As a result, TR-PBO membranes from AcPIs, i.e., cPBO and sPBO, demonstrated over 5000 Barrers of CO2 permeability as shown in Table 1. In comparison of tPBO and aPBO, even though they have the same chemical structure of HPI and resultant TR-PBO, they demonstrated different gas transport behaviors. In the case of tHPI, intermolecular crosslinking is formed during thermal imidization, whereas aHPI retains a linear chain structure during azeotropic imidization. This behavior leads to different physical properties of HPIs and resulting TR-PBOs, such as thermal conversion conditions, free volume elements, cavity sizes, and gas transport properties. Generally, tPBO prepared from tHPI demonstrated larger free volume elements and higher gas permeability compared with aPBO. tPBO demonstrated around 4000 Barrers of CO2 permeability, whereas aPBO yielded 400 Barrers of CO2 permeability when their precursors were thermally rearranged at the same conditions, as shown in Table 1 [121]. The difference in gas transport behaviors mainly results from the free volume elements and cavity size distributions of TR polymers, which will be discussed in the following section. Effect of ortho positioned functional groups in precursor polyimides on physical properties of the resulting TR polymers has been reported [133,134]. Different precursor imidization routes influenced gas separation performances

Table 1 Gas permeability and selectivity of TR polymer and copolymer membranes. Polymer code

CPBO pHAB-6FDA pTR450 mHAB-6FDA mTR450 TR-PBO XTR-PBOI-5 XTR-PBOI-10 XTR-PBOI-15 XTR-PBOI-20 TR-␣-PEBO 450-1 450-3 TR-␣-PBI PBI TR-␤-PBO mPBO pPBO 6fPBO PBO450 TR-␣-PBO-co-PI TR-PBO TR-PBO-co-PI

6FDA + bisAPAF BPDA + bisAPAF ODPA + bisAPAF BTDA + bisAPAF PMDA + bisAPAF 6FDA + bisAPAF 6FDA + bisAPAF 6FDA + bisAPAF 6FDA + bisAPAF 6FDA + HAB-EA 6FDA + HAB-Ac 6FDA + HAB-PAc 6FDA + HAB 6FDA + APAF BTDA + APAF ODPA + APAF 6FDA + HAB 6FDA + HAB(95) + bisAHPF(5) 6FDA + HAB(90) + bisAHPF(10) 6FDA + HAB(85) + bisAHPF(15) 6FDA + bisAHPF (cardo) 6FDA + HAB PI 6FDA + HAB PBO 6FDA + mHAB PI 6FDA + mHAB PBO 6FDA + bisAPAF 6FDA + bisAPAF + DABA(5) 6FDA + bisAPAF + DABA(10) 6FDA + bisAPAF + DABA(15) 6FDA + bisAPAF + DABA(20) 6FDA + 6FBAHPP 6FDA + 6FBAHPP 6FDA + DAB IPCl + bisAPAF TPCl + bisAPAF 6FCl + bisAPAF BPDC + bisAPAF BPDA + bisAPAF BPDA + bisAPAF(8) + ODA(2) BPDA + bisAPAF(5) + ODA(5) BPDA + bisAPAF(2) + ODA(8)

Permeability (Barrers) H2

CO2

2774 444 91 356 635 4194 408 3612 3585

4045 597 73 469 952 4201 398 5568 5903 51 174 211 410 1993 149 112 296 1079 1539 1306 255 10 240 12 720 261 746 980 668 440

530 1665 229 188 407 1189 1479 1254 371 35 260 46 570 294 603 763 515 421 95.3 439

41.4 486

Selectivity O2 747 93 14 81 148 1092 81 1306 1354

N2 156 20 2.3 15 34 284 19 431 350

CH4

Ref.

O2 /N2

CO2 /N2

CO2 /CH4

H2 /CO2

H2 /N2

H2 /CH4

4.8 4.7 6.1 5.4 4.4 3.8 4.3 3.0 3.9

26 30 32 31 28 15 21 13 17

0.69 0.74 1.2 0.76 0.67 1.0 1.0 0.65 0.61

18 22 40 24 19 15 22 8.4 10

38 30 91 36 28 28 34 14 14

4.0 3.1 4.8 4.9 4.4 4.0 3.8 3.8 4.6 6.6 4.5 6.8 3.8 4.2 4.5 3.8 4.0 4.2

16 13 23 21 21 19 18 19 22 29 24 29 21 21 25 19 22 22

55 40 73 47 41 28 34 22 23 36 34 19 23 17 38 35 32 26 23 22 28 63 31 67 23 35 37 30 34 35

100 474 31 26 62.7 227 316 264 54.2 2.3 45 2.8 130 52.5 133 193 119 81.9

25.3 154 6.5 5.3 14.3 57.1 83.6 69.3 11.8 0.35 10 0.41 34 12.6 29.6 50.9 29.8 19.7

73 15 1 10 23 151 12 252 260 1.4 5.1 11.4 18.2 115 3.9 3.2 9.2 41.7 65 58.7 9.2 0.16 7.7 0.18 31 7.5 19.9 33.0 19.4 12.4

10.0 88.5

1.89 20

1.45 17

5.3 4.4

22 24

29 29

337

62.0

35

5.4

26

1779

1624

65 128 65 526

22 72 44 532

6.4 17 11 105

0.4 3.2 5.6 30.3

0.5 1.9 1.5 28.9

1228 623 47 38

1014 389 25 11

220 90 4.8 2.2

48 18 0.82 0.4

41 14 0.65 0.3

16 5.3 2.0 3.5 4.6 5.0 5.9 5.5

[51,122]

[121]

[129]

1.3 0.8 1.5 1.7 1.4 1.1 0.96 0.96 1.5 3.5 1.1 3.8 0.8 1.1 0.8 0.8 0.8 1.0

21 11 35 35 29 21 18 18 31 100 26 112 17 23 20 15 17 21

29 14 59 59 44 29 23 21 40 219 34 256 18 39 30 23 27 34

[127] [134]

2.3 0.9

66 26

66 26

[124]

46

1.1

29

51

[123]

55 23 7.9 18

44 38 29 18

3.0 1.8 1.5 0.99

163 40 12 17

130 67 43 18

[135]

21 22 31 28

25 28 39 37

1.2 1.6 1.9 3.5

26 35 57 95

30 45 72 127

[148]

[126]

[131]

[132]

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

TR-␣-PBO TR-1 TR-2 TR-3 TR-4 TR-5 tPBO aPBO cPBO sPBO TR400 TR400 TR400 TR 450 APAF-6FDA APAF-BTDA APAF-ODPA PBO CPBOc

Structure

[136]

9

10

Table 1 (Continued) Polymer code

Structure

Permeability (Barrers) H2

TR-PPL TR-␤-PBO-co-PA PHBOA(8:2) PBOA(8:2) PHBOA(2:8) PBOA(2:8) TR-␣,␤-PBO PIBO300 PBO400 Thermally labile PI PI-g-CD 400 PI-g-CD 425 PI-g-CD 450 PI-S1-400 PI-S2-400 PI-S3-400 PI-S1-425 PPM-␣-CD-425 PPM-␤-CD-425 PPM-␥-CD-425

BPDA + ODA ODPA + bisAPAF(8) + MDA(2) ODPA + bisAPAF(8) + DAM(2) ODPA + bisAPAF(8) + OT(2) ODPA + bisAPAF(8) + BAP(2) ODPA + bisAPAF(8) + BAPP(2) ODPA + bisAPAF 6FDA + HAB(3) + 4MPD(1) 6FDA + HAB(1) + 4MPD(1) 6FDA + HAB(1) + 4MPD(3) 6FDA + HAB(3) + FDA(1) 6FDA + HAB(1) + FDA(1) 6FDA + HAB(1) + FDA(3) 6FDA + bisAPAF 6FDA + bisAPAF(8) + DAB(2) 6FDA + bisAPAF(5) + DAB(5) 6FDA + bisAPAF(2) + DAB(8) 6FDA + DAB IPCl + HAB(8) + ODA(2) IPCl + HAB(8) + ODA(2) IPCl + HAB(2) + ODA(8) IPCl + HAB(2) + ODA(8) TAC + bisAPAF TAC + bisAPAF 6FDA + DABA + ␤CD 6FDA + DABA + ␤CD 6FDA + DABA + ␤CD 6FDA + DABA + glucose 6FDA + DABA + sucrose 6FDA + DABA + raffinose 6FDA + DABA + glucose 6FDA + durene(9) + DABA(1) + ␣CD 6FDA + durene(9) + DABA(1) + ␤CD 6FDA + durene(9) + DABA(1) + ␥CD

14 43.4 53.2 47.3 31.2 40.4 44.9 – – – – – – 4194 1989 2895 1680 376 1.8 3.42 2.76 4.6 104 663

Selectivity O2

2.7 18.0 23.5 16.8 11.9 18.8 15.7 19 52 226 14 23 36 4201 1874 1805 525 234 0.22 0.64 0.36 0.68 30.2 456 921 4016 8000 533 370 407 1389 2423 3112 4211

N2 0.72 3.90 4.99 3.45 2.39 3.48 3.95 – – – – – –

1092 421 475 132 65 0.07 0.15 0.1 0.22

CH4

0.09 0.66 0.79 0.57 0.39 0.62 1.25 0.90 2.3 11 0.94 1.4 2.6 284 94 85 18 13 0.009 0.024 0.01 0.025

0.03 0.41 0.43 0.32 0.22 0.41 1.74 0.48 1.2 5.8 0.59 1.04 1.47 151 50 46 6.7 8.1 0.008 0.017 0.007 0.013

Ref.

O2 /N2

CO2 /N2

CO2 /CH4

H2 /CO2

8.0 5.9 6.3 6.1 6.1 5.6 3.2 – – – – – –

30 27 30 30 31 31 13 21 23 21 15 16 14

90 44 55 53 55 46 9.0 40 43 39 24 22 24

5.2 2.4 2.3 2.8 2.6 2.2 2.9 – – – – – –

156 81 67 83 80 65 36 – – – – – –

467 106 124 150 145 98 26 – – – – – –

3.8 4.5 5.6 7.3 5.0

15 20 21 29 18

28 38 39 78 29

1.0 1.1 1.6 3.2 1.6

15 21 34 93 29

28 40 63 250 46

[149]

7.8 6.3 10 8.8

24 27 36 27

28 38 51 52

8.2 5.3 7.7 6.8

200 143 276 184

225 201 394 354

[150]

3.4 1.5

68 28

151 38

[151]

8.93 98.1

1.53 23.8

0.69 17.4

5.8 4.1

20 19

44 26

262 1399 2707 135 88.9 106 254 572.77 754.35 1024.4

78 377 523 33.7 21.7 27.8 66.7 127.67 166.1 231.23

58 247 463 21.5 13.6 19.3 51.5 111.67 140.25 187.66

3.4 3.7 5.2 4.0 4.1 3.8 3.8 4.5 4.5 4.4

12 11 15 16 17 15 21 19 19 18

16 16 17 25 27 22 27 22 22 22

H2 /N2

H2 /CH4 [146]

[147]

[137]

[139]

[138]

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

PI PBO-MDA PBO-DAM PBO-OT PBO-BAP PBO-BAPP PBO HAB:4MPD 3:1 HAB:4MPD 1:1 HAB:4MPD 1:3 HAB:FDA 3:1 HAB:FDA 1:1 HAB:FDA 1:3 TR-␣-PBO-co-PPL TR-PBO TR-PBO-co-PPL

CO2

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

of the resulting TR polymer membranes. Sanders et al. investigated the synthesis routes of various precursors and introduced acetate, propanoate, and pivalate linkages on HAB-6FDA based TR polymers [133]. Polymer density and free volume of precursors were influenced by the structure and size of ortho positioned groups. Polyimides with pivalate groups, containing large functional groups than others, showed higher free volume than that of other polyimides. Acetate and propanoate substituted in ortho positioned groups showed lower thermal stability than that of pivalate groups. Polyimides with these substituent groups exhibited lower thermal rearrangement temperatures than polyimides with pivalate or hydroxyl groups [133]. Other approaches have been investigated to develop TR-PBO membrane by introducing isomeric monomers and crosslinked polymers. Introduction of a simple isomeric monomer in the form of a diamine moiety can change the physical and gas separation proper˜ ties of TR-PBO membranes. Comesana-Gándara et al. introduced 3.3 -diamino-4.4 -dihydroxybiphenyl (mHAB), which is an isomeric monomer of commercial 3.3 dihydroxybenzidine (HAB), with a meta linkage of a diamine moiety [131]. Preparation of mHAB investigated by condensation reaction from 4.4 -dihydroxybiphenyl with an easy synthesis route provided high yield as well as cheap price. The precursor polyimide for TR-PBO was prepared with mHAB and 4,4 -hexafluoroisopropylidene diphthalic anhydride (6FDA) monomers, and their corresponding TR-PBO demonstrated a large improvement in gas permeability where isomeric TR-PBO membrane demonstrated 720 Barrers of CO2 permeability with a CO2 /N2 selectivity of 21 and CO2/ CH4 selectivity of 23, respectively. On the other hand, TR-PBO membranes from commercial HAB and 6FDA exhibited 240 Barrers of CO2 permeability and CO2 /N2 and CO2 /CH4 selectivity of 24 and 31, respectively. Isomeric polyimides also exhibited lower glass transition temperature (Tg ) and thermal rearrangement temperature (TTR ). Crosslinked TR-PBO (XTR-PBO) has been reported to investigate the effect of chemical crosslinking in terms of physical properties, cavity sizes, and gas transport behaviors [132]. Precursor polyimides, which have ortho positioned hydroxyl groups as a thermal rearrangement moiety, as well as pendant carboxylic acid groups in small molar ratio as a crosslinking moiety, have been introduced and crosslinked with 3,5-diamino benzoic acid (DABA) as described in Fig. 3(d). Crosslinked XTR-PBO membranes after azeotropic imidization and thermal rearrangement of polyimides overcame the collapse of the free volume during thermal rearrangement with larger cavity sizes and higher fractional free volume. As a result, a synergistic effect of high permeability and high selectivity appeared to be created in one step, in particular for the gas pair CO2 /CH4 . XTR-PBO membrane with 10% crosslinking demonstrated 980 Barrers of CO2 permeability and CO2 /N2 and CO2 /CH4 selectivity of 19.3 and 29.7, respectively. Among the TR polymers, TR-PBI has been reported by post-thermal treatment of TR-PPL to improve the highly permeable characteristics for small gas molecules and selective molecular sieving properties [123]. The

11

ladder-like pyrrolone structure was reorganized to a benzimidazole structure by alkaline treatment and subsequent thermal rearrangement forming TR-PBI. The CO2 permeability of TR-PBI was demonstrated as approximately 1600 Barrers with a CO2 /CH4 selectivity of 46. Commonly, conventional PBI membranes have been investigated for use in high temperature hydrogen separation processes with outstanding hydrogen selectivity [141]. TR-PBI membranes showed a potential of CO2 separation as well as H2 separation, exhibiting a satisfactory H2 /CO2 separation property at high temperature [123]. One of the great advantages of TR polymer is its processability to any types of membranes for module preparation. For gas separation applications, hollow fiber membranes are the most common types for membrane modules due to their large membrane area per volume and packing density with high gas flux, which can reduce the operation footprint [142–144]. Precursor polyimides are soluble in common organic solvents and can be directly spun into hollow fiber membranes. These precursor hollow fiber membranes were converted into TR hollow fiber membranes by post-thermal treatment as shown in Fig. 6 [145]. After thermal treatment, TR hollow fiber membranes were not soluble anymore due to thermal crosslinking, but they offered high chemical and thermal stability. Laboratory scale TR polymer hollow fiber membranes have demonstrated CO2 permeance of approximately 2000 GPU with a CO2 /N2 selectivity of about 20. The previous values were obtained with an effective membrane skin layer thickness of 1.0–1.5 ␮m, which has the possibility of being reduced to the nm scale upon industrial production with a higher gas flux at least ten fold [145]. 3.2.2. TR polymers derived from polyamides with hydroxyl groups TR polymers from functionalized polyamide were prepared using a mild thermal treatment conversion at low temperature. In the case of the TR polymer from hydroxyl polyamides (HPAs), named TR-␤ polymers, HPA precursors were resulted from polycondensation reaction by the reaction of diacid chlorides and bisaminophenoles in polar aprotic solvents as described in Fig. 3(b) and (c). Then, solution process was used for HPA membrane formation. Thermal rearrangement from HPA was performed at around 350 ◦ C in an inert atmosphere as shown in Fig. 7. Those series of TR-␤ polymer membranes exhibited different aspects compared with TR-␣ polymer membrane, especially in considering gas transport properties. TR-␤ polymer membranes demonstrated high gas selectivity for the small gas pair of H2 and CO2 with promising H2 permeability [135,136]. H2 permeability of TR-␤-PBO membranes was observed from 65 Barrers to around 150 Barrers with a H2 /CO2 selectivity of around three depending on the polymer structure and thermal rearrangement conditions [135]. H2 permeability of TR-␤-PBO membranes was an order of magnitude higher than conventional H2 separation membranes, but offers a low selectivity. Temperature dependence of TR-␤-PBO membranes on H2 and CO2 separation was also investigated, since the actual application for H2 and CO2 separation in an IGCC plant is operated at high temperature and pressure

12

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

Fig. 6. Hollow fiber membranes made of (a) hydroxyl poly(amic acid), (b) hydroxyl polyimide, and (c) TR-PBO [145]. Copyright 2012. Reproduced with permission from Elsevier Ltd.

Fig. 7. Thermal rearrangement mechanism from hydroxyl polyamide to polybenzoxazole (TR-␤-PBO) [135].

Fig. 8. Temperature dependence on H2 permeability and H2 /CO2 selectivity (a) with temperature variation (closed: H2 permeability, open: H2 /CO2 selectivity) and (b) with polymeric upper bounds (: mPBO, 䊉: pPBO, : 6fPBO, : TR-PBI) [123]. Copyright 2010. Reproduced with permission from [135], Copyright 2012. Reproduced with permission from Elsevier Ltd and the Royal Society of Chemistry, respectively.

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

13

Fig. 9. Polyimides with thermally labile units and their labile units: (a) glucose (S1), (b) Sucrose (S2), (c) Raffinose (S3), (d) ␣-cyclodextrin (␣-CD), (e) ␤-cyclodextrin (␤-CD), and (f) ␥-cyclodextrin (␥-CD).

[29]. Gas permeability is easily affected by temperature due to diffusion and sorption phenomena. Basically, as temperature rises, gas diffusion is accelerated and gas solubility is decreased due to the thermal motion of gas molecules. According to the solution–diffusion mechanism, gas permeation and separation properties in H2 and CO2 separation are both improved which, is an unusual case of common membrane separation gas pairs. Permeability of TR-␤-PBO membranes of both H2 and CO2 were increased as a function of operation temperature, however the permeability ratio increase of H2 was much higher than that of CO2 , consequently, H2 /CO2 selectivity also increased. The highest separation performance with reasonable gas permeability and selectivity was 206 Barrers of H2 permeability with a H2 /CO2 selectivity of 6.2, as shown in Fig. 8 [135]. 3.2.3. Polyimides with thermally labile units Modified polyimides with thermally labile units brought out rapid gas transport and reduced plasticization properties. Such polyimides were synthesized from dianhydrides and diamines with functional groups, including sulfonated groups ( SO3 H) and carboxylic acid groups ( COOH), which can be substituted with various large labile units, such as glucose, sucrose, raffinose, and various cyclodextrins (␣,␤,␥,-CDs), as shown in Fig. 9. Those polymers were prepared into polymer membrane films

and the thermally labile units in polymer structure were decomposed by post-thermal treatment at around 400 ◦ C. Thermal decomposition of large functional groups in a solid state increased free volume elements and improved membrane gas permeability. However, the treatment temperature must carefully be selected to prevent pyrolysis [137–140]. The highest improvement of CO2 permeability is from PI grafted with ␤-CD. The CO2 permeability of those PI membranes after decomposition of thermally labile units demonstrated 8000 Barrers with a CO2 /CH4 selectivity of 17 at 35 ◦ C [137]. 3.2.4. TR co-polymers Various TR copolymers were also investigated due to the synergy effects of various TR polymers and glassy polymers [146–151]. TR poly(benzoxazole-co-imide) (TR-PBO-coPI) [146–148] and TR poly(benzoxazole-co-pyrrolone) (TR-PBO-co-PPL) [149] were thermally rearranged from copolymer precursors, which were reported from a dianhydride monomer with two different diamine monomers. TR poly(benzoxazole-co-amide) (TR-PBOA) membrane was also prepared from copolymer precursors synthesized from an acid chloride monomers with two different diamine monomers as described in Fig. 3 [150]. TR-␣,␤ copolymers have also been investigated to achieve a synergetic effect between TR-␣ polymers and TR-␤ polymers with good permeability and selectivity by thermal rearrangement of

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S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

poly(hydroxyamide amic acid). However, flexible imide linkages in polymer structure reduced their gas separation properties, but improved the membrane processability [151]. The gas transport performances of TR-PBO-co-PI membranes were easily controlled by varying the monomer copolymer ratio during the imide step, as can be seen in Table 1. Cavity sizes and size distributions created during the thermal rearrangement process could be easily controlled with various functional polyimide copolymer compositions. TR-PBO-co-PI membranes prepared from 2,2-bis(3-Amino-4-hydroxyphenyl)hexafluoropropane (bisAPAF), 4,4-oxydianiline (ODA), and 3,3,4,4-biphenyltetracarboxylic dianhydride (BPDA) exhibited the CO2 permeability ranged from 2.7 Barrers to 1014 Barrers depending on the membrane composition without a significant loss in selectivity. The flexible PI part in TR-PBO-co-PI membranes can also improve the mechanical properties of TR polymer membranes [148]. However, TR-PBO-co-PI membranes prepared from 4,4 -oxydiphthalic anhydride (ODPA) and bisAPAF with various diamines [146], or 6FDA, HAB, and two different diamines, i.e., 2,3,5,6-tetramethyl-1,4-phenylenediamine (4MPD) or 9,9 -bis (4-aminophenyl)fluorene (FDA) [147] demonstrated relatively lower CO2 permeability, listed in Table 1. TR-PBO-co-PPL was introduced to improve the gas selectivity. PPL part in the copolymer enhanced gas selectivity due to its well-packed ladder-like structure that improved interchain packing and the molecular sieving effect of TR copolymer membrane. TR-PBO-co-PPL was prepared from precursors of polyimides containing both hydroxyl and amino groups, following the thermal conversion into TR-PBO and TR-PPL. TR-PBO-co-PPL membranes in various copolymerization ratios, exhibited higher gas permeability and selectivity than those of their corresponding precursor polymers. Moreover, in terms of gas selectivity, copolymerization of TR-PBO-co-PPL showed synergetic effects with enhanced gas selectivity, yet without losing significant gas permeability [149]. TR-PBOA copolymer membranes were reported from poly(o-hydroxyamide-co-amide) (PHAA) precursors following in situ thermal cyclization reaction [150]. These copolymer membranes demonstrated very high gas selectivity especially for H2 separation properties. Precursor PHAA contained both hydroxyamide part from HAB moiety and amide part from ODA moiety. Rigid biphenyl group from HAB moiety in polymer structure increased gas selectivity, whereas flexible ether group from ODA moiety improved membrane processability and gas permeability. Gas permeability and selectivity of the resulting TR copolymer membranes, which were varied depending on the ratio of hydroxyamide and amide parts, achieved H2 /CO2 selectivity up to 8.4 with H2 permeability of 1.8 Barrers at 35 ◦ C and 10 bar [150]. Moreover, both H2 permeability and H2 /CO2 selectivity of their copolymer membranes significantly increased at high operating temperature because of the difference in activation energy of permeation (Ep ) between H2 and CO2 . Poly(hydroxylamide amic acid) precursor was introduced to investigate more uniform TR-PBO copolymers,

so-called TR-␣,␤ copolymers [151]. Those TR-␣,␤ copolymers have the same final TR-PBO structures. However, they demonstrated different properties due to the different thermal rearrangement mechanism routes from HPI for the TR-␣ polymer part and HPA for the TR-␤ polymer part, respectively. The ratio of HPI and HPA was varied to observe the effect of each part, and then the precursors were thermally rearranged into TR-␣,␤ copolymers. Those TR␣,␤ copolymer membranes demonstrated gas permeability and selectivity between respective TR polymer membranes of TR-␣-PBO and TR-␤-PBO. The results explain that the gas transport performances of TR polymer membranes can be well-tuned by appropriate process conditions. Highly permeable TR-␣-PBO membranes and highly selective TR␤-PBO led to high permeability or high selectivity when TR-␣,␤ copolymer membranes were tested. CO2 /CH4 selectivity reached 49.4 with CO2 permeability of 5.06 Barrers when TR-␤-PBO portion was dominant, while CO2 permeability reached 456 Barrers with CO2 /CH4 selectivity of 26.2 when TR-␣-PBO was dominant. TR-␣,␤ copolymer membranes are still attractive as they show high gas permeability and selectivity [151]. Various TR polymers and their hollow fiber membranes have been reported for gas separation applications. TR polymer membranes usually exhibit extraordinary gas permeability for small gases compared to other polymer membranes. Table 1 summarizes gas permeation and separation properties of TR polymer membranes from the previous sections. Gas permeation properties of TR polymer were dependent on the chemical structure of precursor and resulting TR polymer membranes. Structure–property relationship has been reported for gas permeation properties of polyimide membranes [52]. Polymer packing density and chain mobility influence gas permeability and selectivity of polyimide membranes. Therefore, polyimides with bulky 6F group (hexafluorine group) showed high gas permeability with low packing density [52]. From Table 1, similar trends in gas transport properties of TR polymers were observed; 6F group containing TR polymer membranes usually exhibited high gas permeability. However, various factors including precursor imidization method or thermal rearrangement mechanism also strongly influences as well. 3.2.5. Physical property and characterization of TR polymers Various physical properties have been studied to characterize intrinsic polymer properties such as sorption properties, cavity size analysis, and thermal relaxation characteristics [53,152–165]. Both experimental and simulation methods were introduced to characterized TR polymers. Those analysis methods have been widely used to characterize the polymeric materials, especially for membrane materials [52]. Previously, such methods have been applied for the characterization of TR-␣ polymer. However, they can also be used for other TR polymer series. Sorption properties of small gas penetrants have been studied to characterize individual gas diffusivity and solubility for the gas permeability by gravimetric and volumetric methods for small gas molecules, as well as inverse gas chromatography (IGC) for large gas molecules. Study

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

15

Fig. 10. Cavity size analysis of (a) precursor polyimides and (b) TR-polybenzoxazoles by positron annihilation lifetime spectroscopy (PALS) [121]. Copyright 2010. Reproduced with permission from the American Chemical Society.

on gas sorption and permeation were usually performed to elucidate the gas transport mechanism of polymer membranes [166]. Thermal rearrangement enhances the free volume, which especially leads to diffusion enhancement with only a small increase in sorption for small molecule transport. Small gas molecules permeate through the polymer matrix, so-called free volume element according to the solution–diffusion mechanism, which is described by the diffusivity and solubility coefficient relationship. Sorption isotherms of TR polymers and copolymers showed the sorption behavior of typical glassy polymers, which are concave to the pressure axis so-called dual-mode sorption. Their sorption coefficients exhibited higher values than that of typical glassy polymers. Sorption properties of TR polymers also explained that thermal rearrangement increased the solubility coefficient of the precursor polymer by enhancing the free volume element [153,163,165]. However, it has been concluded that the enhancement of gas permeability for small gases by thermal rearrangement is mainly influenced by an increase in gas diffusivity coefficient, where the increment of sorption coefficient is not critical. For the sorption of large gas molecules such as hydrocarbons, thermal rearrangement resulted in a great increase in the solubility coefficients. Moreover, TR polymers showed strongly negative mixing parameters during IGC measurement, indicating the large free volume in TR polymers, which played an important role in sorption of large gas molecules [155]. Cavity sizes and microporous characteristics of polymers have been analyzed using positron annihilation lifetime spectroscopy (PALS) and Brunauer–Emmett–Teller (BET) analysis. PALS analysis provides the free volume elements in polymers from the annihilation lifetime of ortho positron (oPs). PALS measurement is appropriate to investigate cavity sizes and size distributions to explain the gas transport mechanism through polymeric membranes [167]. Generally, the cavity sizes in polymeric membranes exhibit a broad distribution that disturbs efficient gas separation, where precursor polyimides also demonstrate broad cavity size distribution. However, TR polymers represented bimodal cavity distributions where the large

cavities are efficient for rapid gas diffusion, while the small cavities are appropriate for gas separation by size difference in gas molecules as shown in Fig. 10. In TR-␣ polymer, the large cavity sizes are about 8.0–10 A˚ and the small sizes ˚ depending on the type of TR polyare about 3.0–4.0 A, mers. The large cavities contribute to rapid gas transport, whereas the smaller cavities are suitable for CO2 separation [109,128]. According to Fig. 10, the difference in cavity sizes among TR polymers, i.e., tPBO, aPBO, cPBO, and sPBO, exhibited up to approximately a ten-fold difference in CO2 permeability, as shown in Table 1, even though the difference in cavity size is less than 10%. Those well-tuned cavity ˚ which is slightly larger than the kinetic sizes less than 3.8 A, diameter, suggested by Breck and commonly used for gas transport through membranes [168], of CO2 molecules ˚ yet similar to that of CH4 (3.8 A) ˚ and N2 (3.64 A), ˚ (3.3 A), led to the rapid gas transport of CO2 and good selectivity over CH4 and N2 [168,169]. BET analysis is commonly used to analyze microporous materials by typically measuring nitrogen adsorption/desorption isotherms at low temperature. Fig. 11 shows nitrogen adsorption/desorption isotherms of TR polymers measured by BET analysis measured at −195 ◦ C [51]. The nitrogen adsorption/desorption isotherms of TR-1 and TR-2 showed irreversible Type 1 form with hysteresis. The BET surface areas of TR polymers were remarkably large; TR-1 showed 510 m2 /g, which indicated the presence of large free volume [51]. The hysteresis for the TR polymers was similar to that of PIMs [170]. The relationships between the polymer structure and cavity sizes, as well as the transport properties, have been studied by computer simulation approach [53,157,162,164]. Optimized simulation procedures revealed the gas diffusion through polymer matrices with different cavity sizes and shapes. Moreover, molecular dynamics (MD) simulation on a nanosecond (ns) scale can characterize structural properties of TR polymers. This simulation study demonstrated the property change from precursor polymers to TR polymers during thermal rearrangement in terms of free volume elements (Fig. 12), and angle distribution [157]. Gas permeability from diffusivity and solubility was also calculated from

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S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

Fig. 11. Nitrogen adsorption/desorption isotherms at −195 ◦ C for (a) precursor of TR-1, (b) precursor of TR-2, (c) TR-1, and (d) TR-2. p/p0 is the ratio of gas pressure (p) to saturation pressure (p0), with p0 = 746 Torr [51]. Copyright 2013. Reproduced with permission from AAAS.

MD simulation method. Gas solubility and diffusivity of TR polymers from constructed polymer model by MD simulation showed good correlation between simulated and experimental results [162,164]. Cavity sizes and size distributions in polymer membranes directly affect gas diffusion through membranes. According to Thornton et al., gas diffusion is considered via activation, surface, and Knudsen diffusion by the cavity sizes, where diffusion and sorption by the solution–diffusion mechanism have different contributions in each case [52,171]. These different diffusion mechanisms dominate transport of small gas molecules and are determined by pore sizes. The smallest pores where gas molecules attempt to enter dominates activation diffusion to make a diffusive jump with molecular sieving

effect. Surface diffusion is dominated in the pore size region between activation where the surface area of the pore wall increases for more gas molecules adsorption on the wall. A diffusive jump in surface diffusion occurs from one adsorption site to the other. Knudsen diffusion is known to apply ˚ In this region, the to pores size between 10 A˚ and 500 A. permeation rates of gas molecules are dependent on the gas molecular weight [52,171]. From the mathematical formulation, suction energy was considered as a function of cavity sizes and gas penetrants, and it determines the kind of diffusion mechanism by cavity sizes [171]. This approach ˚ provided that in the cavity sizes between 3.0 A˚ and 10 A, which was similar to the cavity sizes of TR polymers, diffusion is dominated by the suction energy, that is surface diffusion [171]. In the case of TR polymers, their small cavity sizes support activation diffusion and surface diffusion, which improved gas diffusivity as well as diffusion selectivity. The thermal relaxation of TR polymer has been studied using dynamic mechanical analysis (DMA) to characterize the thermal motion of polymers in ␣ (the glass-rubber), ␤, and ␥ (sub-glass) relaxations [154]. In thermal behavior of polymers, ␣ relaxation explains general glass-rubber transition where rigid and glassy polymer chain becomes relaxed and rubbery state by ␣ relaxation. Sub-glass relaxations, which are ␤ and ␥ relaxations, explain movement of side chain and local mobility in non-equilibrium state, respectively, below Tg . Sub-glass relaxations are from chain stiffness or rigidity [154]. Various HPI precursors were prepared and series of dynamic mechanical studies were performed for both precursors, and TR polymers prepared by chemical and thermal imidization. The effect of the thermal rearrangement process as a function of polymer structure and prior thermal exposure history for the thermal conversion was observed by the change of the glass transition, i.e., ␣ transition. For precursor polyimides, two tan ı curves were presented as a function of temperature, where the first and second curves represented Tg and TTR , respectively. Storage modulus and tan ı

Fig. 12. Free volume distribution (blue color) of (a) HPI and (b) TR-PBO [157]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.) Copyright 2012. Reproduced with permission from the American Chemical Society.

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

17

Fig. 13. Synthesis of PIMs and some examples of PIMs (a) PIM-1 and (b) PIM-7; (synthesis reagents and conditions, i: Monomer A, Monomer B, K2 CO3 , DMF).

curves were affected by chemical structure of the precursor polymers, as well as imidization methods. Dynamic mechanical characteristics also provide molecular information in the sub-glass transition region regarding the polymer mobility and local relaxation. For the ␤ transition, thermal rearrangement decreased the relaxation intensity and shifted the transition temperature to lower temperature (80–100 ◦ C), meaning localized chain motions from a compact interchain distance, as well as the enhancement of free volume underlying the ␤ response. The ␥ transition process was absent, which suggests the presence of intramolecular and intermolecular interaction across the polymer matrix. The absence of the ␥ transition explains the strong interaction by both thermal crosslinking during thermal rearrangement and hydrogen bonding of HPI and AcPI [154]. Thermal rearrangement conditions have been studied by thermal analysis, such as TGA and differential scanning calorimetry (DSC), to explore the relationship between chain mobility, glass transition, and thermal rearrangement [130,159–161]. From the analysis of thermal conversion mechanism of precursors, TTR is related with the chain mobility of precursor polymers, which is characterized by the glass transition, where precursor polymers with high chain flexibility demonstrate a low TTR [130,161]. The gel fraction of TR polymers provides information on the extent of intermolecular rearrangement or cross-linking. The gel fraction is the weight percent of insoluble portion of the membrane. After thermal rearrangement at a certain temperature, the gel fraction of treated polymers was increased and after complete conversion, the gel fraction reached 100%. Consequently, thermal rearrangement causes intermolecular and intramolecular reactions, which form insoluble and cross-linked networks [130]. Chain rigidity of precursors also affects the thermal conversion temperature, which should be considered when preparing TR polymers. Thickness dependence of TR

membranes and their physical aging was also investigated with 50 nm to 100 ␮m thick membrane films prepared by either spin coating or knife casting. Membrane thickness increase during thermal rearrangement was observed and thickness increase in thinner film was more severe. Moreover, thermal rearrangement temperature of thinner film was milder comparing with thick film [160]. Thin TR films experienced an intensive physical aging as indicated by rapid gas permeability decrease over time [172]. The atmospheric effect on thermal rearrangement was also important because of oxidation effect at thermal rearrangement temperature. The effects of purge atmosphere were investigated during thermal rearrangement at 425 ◦ C under nitrogen and air atmosphere [159]. Consequently, TR polymer membranes prepared under air exhibited higher gas permeability, but deteriorated mechanical properties. 3.3. Polymers of intrinsic microporosity Polymers with a rigid ladder-like backbone and contorted structure can achieve high free volume elements from restricted chain rotation in the polymer matrix. Polymers of intrinsic microporosity (PIMs) are developed as microporous network polymers with extremely stiff ladder-like domains with a flexible dioxane structure that demonstrate well-defined microcavities [10,173,174]. The intrinsic microporosity was defined as a continuous network of interconnected intermolecular microcavities, which is formed as a consequence of the shape and rigidity of the polymer structure [175]. Since Budd and McKeown first reported new ladderlike polymers with high surface area incorporating catalytic centers in 2002 [176,177], several properties of these interesting PIMs have been studied including microporosity, sorption, and diffusion properties [10,113,114,170,178]. The microporosity possessed in PIMs demonstrated a BET surface area of 500–1730 m2 /g by nitrogen adsorption

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S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

Fig. 14. The nitrogen adsorption/desorption isotherm for a powder sample of PIM-7 [180]. Copyright 2008, the American Chemical Society.

[170,179,180]. One of the first reported PIMs, PIM-7 has a BET surface area of 680 m2 /g characterized from adsorption/desorption isotherm curve as shown in Fig. 14 [180]. The adsorption/desorption isotherm of PIM-7 displays microporous characteristics from similar uptakes with conventional microporous materials such as zeolites. It also

indicates the presence of some mesoporosity and the hysteresis at low pressures resulting in the accessibility of mesopores and micropores [180]. High surface area and solution-processability of PIMs widen their various applications, including hydrogen storage and membrane gas separation [115,173,175,179]. The first reported PIMs series are prepared by a polycondensation reaction of relatively stiff tetrahydroxyl-monomers and tetrafluoromonomers (or tetrachloro-monomers) with contorted centers, as shown in Fig. 13. Monomers to synthesize the PIMs are usually synthetic, so the variation of PIM structures is extended using any number of different bondforming reactions [181]. Contortions in monomers are generally presented by spiro centers (i.e., a tetrahedral carbon shared by two rings) or other rigid non-planar structural units [175]. Because microcavities of PIMs were developed from stiff structure, novel PIMs, such as PIMpolyimide or PIMs with more rigid structure based on Tröger’s base (TB) unit or ethanoanthracene (EA) formation, have recently been reported [116,182]. Fig. 15 shows examples of monomers used for the synthesis of PIMs. Solution-processability of PIMs depends on the choice of monomers where only some PIMs form a network polymer without solvent solubility. Soluble PIMs such as PIM-1, prepared from monomers of A1 and B4, and PIM7, prepared from monomers of A1 and B8 in Fig. 15, are processed to self-standing films or coated films by simple casting from polymer solution, and are transparent and

Fig. 15. Monomers A and B for synthesis of PIMs.

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

19

Fig. 16. Flexible and transparent PIM-1 film [183]. Copyright 2008. Reproduced with permission from the American Chemical Society.

flexible as shown in Fig. 16 [10,183]. Thermal and chemical stability of PIMs have been investigated at high temperature and under harsh chemical conditions. PIM-1 revealed no glass transition and melting temperature measured by dynamic mechanical analysis [179]. They also showed great stability in acidic, basic, and oxidation conditions [175]. The contorted and rigid backbones of PIMs restrict the rotational free energy and prevent effective packing with a robust mechanical strength. They also provide anti-plasticization properties [109,184]. The intrinsic microporosity of PIMs is generated by a rigid structure that retains efficient pore sizes of free volume elements for rapid gas transport by high diffusion coefficients. Micro˚ as measured by PALS, cavities in the range of 6.0–8.0 A, are efficient to represent a high gas permeation property by a surface diffusion model [184]. Soluble PIMs, such as PIM-1 and PIM-7, show a substantially high selectivity with extraordinary permeability, which represents an advance across the upper-bound for O2 /N2 separation. Their molecular structures satisfied the chain rigidity requirement and the prohibition of chain rotation was achieved by fused ring structures, which could improve gas permeability and selectivity. Reported O2 permeability of PIM-1 were ranged from 370 Barrers to 2270 Barrers, while the CO2 permeability were ranged from 2300 Barrers to 13,600 Barrers, as summarized in Table 2 [184]. Gas transport properties of PIM membranes varied depending on polymer preparation parameters, such as a choice of solvent, monomer purities, and methanol post-treatment. In particular, methanol treatment during or after membrane formation, before gas permeability measurement, improved gas permeability nearly three-fold compared with membranes without methanol treatment [185,186]. PIMs after methanol treatment experiences intensive physical aging. Oxygen permeability was reduced by about 23% and nitrogen permeability by about 40% after 45 days [186]. Gas and vapor sorption of PIM-1 has been studied with experimental and modeling methods [187–189]. Sorption isotherms of PIM-1 for small gases such as CO2 were described by the dual mode sorption, which is similar to that of TR polymers. PIM-1 and PIM-7 also demonstrated a high CO2 permeability with high selectivity of CO2 /N2 and CO2 /CH4 , which surpassed the upper-bound [170]. The membrane

preparation method, such as methanol swelling and water contact, would change the gas permeability of PIM membranes [185,186]. Most of the gas permeability values reported in the literatures are from methanol-treated PIMs. Various structural modifications have been investigated to improve gas transport properties of gas separation membranes by disrupting the efficient space packing for intrinsic microporosity [117,180,190]. 3.3.1. Ethanoanthracene (EA) and Tröger’s base (TB) based PIMs Note that the high gas transport of PIM membranes resulted from the microporosity provided by intrinsic chain rigidity. A rigid spirobisindane (SBI) unit in PIM-1 presented a high gas transport of PIM-1 membranes. However, the relatively flexible dioxane linkage hindered overall rigidity. Recent studies on PIMs as gas separation membranes have been conducted to improve their rigidity by introducing more inflexible units, such as ethanoanthracene (EA) and Tröger’s base (TB), as shown in Fig. 17 [108,116]. A new series of rigid PIMs, using EA and TB units (PIM-EA-TB), as well as PIMs with SBI and TB units (PIMSBI-TB), have been prepared for gas separation membranes. The relative flexibility of the contorted spiro-centers and dioxane rings in PIM-1 was restricted by the new polymer design. Although PIM-EA-TB does not have a contorted part originating from spiro-centers, the combined structure of EA and TB units provides PIM-like fused-ring structures with an intrinsic microporosity, which is described in Fig. 18. Note that the CO2 permeability of PIM-SBI-TB membranes is higher than that of H2 due to the sorption site in the TB unit. However, PIM-EA-TB membranes have demonstrated outstanding molecular sieve properties from entire rigid chain structure in EA and TB units, where the order of gas permeabilities is H2 > CO2 > He > O2 > CH4 > N2 , which is different from the order observed in PIM and TR polymer membranes. Moreover, PIM-EA-TB membranes exhibited a remarkable gas selectivity of O2 over N2 , and H2 over other gases attributed mainly to the diffusivity selectivity for molecules with smaller kinetic diameters. The O2 permeability showed 900–1100 Barrers with an O2 /N2 selectivity of 4.5–6.1 [116]. Furthermore, amines in the TB units would offer simple cross-linking reactions or functionalization to tailor membrane properties for diverse gas

20

Table 2 Gas permeability and selectivity of PIM membranes. Polymer code

Structure

Permeability (Barrers) H2

Basic PIMs PIM-1

A1+B4

Ref.

O2

N2

CH4

O2 /N2

CO2 /N2

CO2 /CH4

H2 /CO2

H2 /N2

H2 /CH4

1300 2332 1900 5010 3580 320 1100 2347 6320

2300 3496 4030 13,600 8310 430 860 4646 13,900

370 786 990 2270 1790 59 190 907 2640

92 238 270 823 727 13 42 242 786

125 360 350 1360 – 22 62 – 1100

4.0 3.3 3.7 2.8 2.5 4.6 4.5 3.7 3.4

25 15 15 17 11 33 21 19 18

18 9.7 12 10 – 20 14 – 13

0.57 0.67 0.47 0.37 0.4 0.74 0.13 0.51 0.45

14 9.8 7.0 6.1 4.9 25 2.6 9.7 8.0

10 6.5 5.4 3.7 – 15 1.8 – 5.7

[170] [184] [183] [117] [195] [180] [170] [190] [117]

7760 2200

7140 2900

2150 720

525 232

699 450

4.1 3.1

14 13

10 6.4

1.1 0.76

15 9.5

11 4.9

[116]

2221 3872 2818 2247 – – – – –

3083 4000 1869 724 2345 1987 1996 1536 1291

483 582 416 189 554 411 375 238 231

101 96 73.6 27.7 161 118 103 57.6 49.9

91 73 62.1 23.1 192 157 116 86.4 52.6

4.8 6.1 5.7 6.8 3.4 3.5 3.6 4.1 4.6

31 42 25 26 15 17 19 27 26

34 55 30 31 12 13 17 18 25

0.72 0.97 1.5 3.1 – – – – –

22 40 38 81 – – – – –

24 53 45 97 – – – – –

[185]

– – – – – – – – – – 2379 1521 92 610

731 1476 2841 3616 1408 1077 2154 1391 2509 3076 4586 2656 150 1120

156 308 561 737 322 216 369 269.6 – – 1053 540 19 140

33 75 158 217 88 52 93 62.7 87 101 369 164 3.9 37

– – – – – – – – – –

22 20 18 17 16 21 23 22.2 29 31 12 16 38 30

– – – – – – – – – – – – 17 20

– – – – – – – – – – 0.5 0.6 0.6 0.5

– – – – – – – – – – 6.4 9.3 24 16

– – – – – – – – – – – – 11 11

[194]

8.7 56

4.7 4.1 3.6 3.4 3.7 4.2 4.0 4.3 – – 2.9 3.3 4.9 3.8

– – – – 2785 395 1015 3670

17 56 558 1953 4572 41 341 6200

3.4 11 116 400 930 32 136 1480

0.57 2 23.3 97.7 269 4.3 23 460

0.5 1.8 22.3 122 374 3.4 22 –

6.1 5.8 5.0 4.0 3.5 7.4 5.9 3.2

30 28 24 20 17 9.5 15 15

34 31 25 16 12 12 16 –

– – – – 0.6 9.6 3.0 0.59

– – – – 10 92 44 8.0

– – – – 7.4 116 46 –

[201]

[192] [193]

[214]

[196] [118] [195] [198]

[200]

[203]

S. Kim, Y.M. Lee / Progress in Polymer Science 43 (2015) 1–32

Cardo-PIM-1 A8+B8 PIM-7 A1+B8 DNPIM-33 A1+A3+B4 PIM-SBF A9+B4 EA and TB based PIMs PIM-EA-TB Fig. 17 PIM-SBI-TB Cross-linked PIMs PIM-300-1.0d Fig. 19 PIM-300-2.0d PIM-UV 20 min PIM-UV 30 min DC-PIM1 DC-PIM2 DC-PIM3 DC-PIM4 DC-PIM5 PIMs with substituted pendant groups TFMPS-PIM1 A1+B12 TFMPS-PIM2 A1+B12(2)+B4(1) A1+B12(1)+B4(1) TFMPS-PIM3 TFMPS-PIM4 A1+B12(1)+B4(2) A1+B11(1)+B4(2) DS-PIM1-33 A1+B13(1)+B4(2) DS-PIM2-33 A1+B10(1)+B4(2) DS-PIM3-33 MTZ100-PIM A1+B17 TZPIM-1 A1+B4+B15+B16 TZPIM-2 A1+B4+B15+B16 ◦ Carboxylated PIM-1 (65 C 1h) Fig. 20 ◦ Carboxylated PIM-1 (65 C 8h) Thioamide-PIM-1 Fig. 20 Thioamide-PIM-1 (EtOH treated) PIM blends and mixed matrix membranes PIM-1/Matrimid (10:90) PIM-1(1)+Matrimid(9) PIM-1/Matrimid (30:70) PIM-1(3)+Matrimid(7) PIM-1/Matrimid (70:30) PIM-1(7)+Matrimid(3) PIM-1(9)+Matrimid(1) PIM-1/Matrimid (90:10) PIM-1(9)+Matrimid(1)+EDA 2hr EDA 2hr TETA PIM-1(9)+Matrimid(1)+TETA PIM-1(9)+Matrimid(1)+BuDA 2hr BuDA PIM-1+Si(13 vol%) Si 13%

CO2

Selectivity

PIM-1+Si(19.1 vol%) PIM-1+Si(23.5 vol%) PIM-1+SWCNT(1 wt%) PIM-1+MWCNT(1 wt%) PIM-1+MWCNT(2 wt%) PIM-1+ZIF-8(28 vol%) PIM-1+ZIF-8(36 vol%) PIM-1+ZIF-8(48 vol%) PIM-1+CC3 (10 wt%) PIM-1+CC3 (30 wt%) PIM-1+CC3 (30 wt%), EtOH treated PIM-1+Silicate-1 (35.5 vol%)

5060 6362 – – – 2980 5745 6680 – – – 894

10,100 12,182 15,721 7813 12,274 4270 6820 6300 3250 5430 37,400 2530

Fig. 21

530 220 360 300 350 1600 1020 840 670 625 259 190

1100 210 520 420 510 3700 2270 2180 2154 1523 263 198

429 261 –

675 263 635

Fig. 22

Fig. 22

2330 3221 2305 1054 1680 870 1640 1680 560 820 6810 351 150 39 85 64 77 545 320 295 270 208 45.2 30.5 120 48 –

880 1573 949 417 713 195 380 350 150 270 3270 83

– – 1820 785 1483 230 510 430 230 480 7220 183

2.7 2.1 2.4 2.5 2.4 4.5 4.3 4.8 3.7 3.0 2.1 4.2

12 7.8 17 19 17 22 18 18 22 20 11 30

– – 8.6 10 8.3 19 13 15 14 11 5.2 14

0.50 0.52 – – – 0.7 0.8 1.1 – – – 0.4

5.8 4.0 – – – 15 15 19 – – – 11

– – – – – 13 11 16 – – – 4.9

[204]

[205]

[207]

[206]

47 9 23 16 19 160 100 94 84 65 10.8 6.9

77 9 27 20 27 260 170 170 168 129 9.1 7.7

3.2 4.3 3.7 4.0 4.1 3.4 3.2 3.1 3.2 3.2 4.2 4.4

23 23 23 26 27 23 23 23 26 23 24 29

14 23 19 21 19 14 13 13 13 12 29 26

0.48 1.1 0.69 0.71 0.69 0.43 0.45 0.39 0.31 0.41 0.98 0.96

11 24 16 18 17 10 10 8.9 8.0 9.6 24 28

6.9 24 13 15 13 6.2 6.0 4.9 4.0 4.9 29 28

[182,208]

30 11 –

34 15 –

4.0 4.4 –

23 24 –

20 18 13

0.64 0.99 –

14 24 –

13 18 –

[212]

[210]

[209]

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Si 19.1% Si 23.5% 1 wt% SWCNT 1 wt% f-MWCNT 2 wt% f-MWCNT ZIF-8 28% ZIF-8 36% ZIF-8 48% PIM-1/CC3-10 PIM-1/CC3-30 PIM-1/CC3-30Et PIM-MFI3 PI-PIM PIM-PI-1 PIM-PI-2 PIM-PI-3 PIM-PI-4 PIM-PI-7 PIM-PI-8i PIM-PI-8ii PIM-PI-9 PIM-PI-10 PIM-PI-11 PIM-6FDA-OH PIM-PMDA-OH PIM-TR-PBO spiroTR-PBO-6F spiroTR-PBO-PM PIM-6FDA-OH

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Fig. 17. The synthesis and molecular structures of EA and TB based PIMs, (a) PIM-EA-TB and (b) PIM-SBI-TB (DMM, dimethoxymethane, TFA, trifluoroacetic acid) [116].

separation applications. Appropriate PIM structure with large free volume achieves enhanced micropore volume where Langmuir affinity mainly contributes to gas sorption. The inclusion of TB group in PIM enhanced Langmuir affinity toward CO2 molecules [189]. However, the sorption of various vapors such as hydrocarbons, alcohols, and water, in TR based PIM showed S-shaped isotherms, indicating a combination of dual mode and Flory–Huggins sorption behavior [187].

3.3.2. Cross-linked PIMs Cross-linking stabilizes the polymer structures by forming a network polymer. It also reduces plasticization produced by condensable gas permeation and improves the gas selectivity [68,191]. Cross-linking methods have been introduced in PIM membranes to improve the gas separation properties using thermal energy and UV irradiation, as shown in Fig. 19 [185,192,193]. Nitrile groups in PIM1 were thermally cross-linked thereby forming the bulky

Fig. 18. The chains of microporous polymers that have a contorted ladder structure resist rotation, which allows them to maintain a more uniform pore structure. (a) The PIM-1 polymer contains SBI and dioxane linkages, but both can bend and flex to a considerable extent. (b) The PIM-EA-TB, which has EA and TB units and has a more rigid architecture [108]. Copyright 2013. Reproduced with permission from AAAS.

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Fig. 19. Cross-linked PIMs: (a) thermally self-cross-linked PIMs, (b) decarboxylation-induced cross-linked PIMs, and (c) UV-rearranged PIMs [185,192,193].

triazine rings [185]. Another method of thermal crosslinking was that after substitution of a nitrile group for carboxylic acid groups, thermal decarboxylation formed a phenyl radical and they were cross-linked by thermal and radical reaction [193]. UV irradiation has been investigated on PIM-1 membrane films. UV irradiation activated homolytic cleavage at the C H bond on the SBI unit and rearranged the polymer structure, resulting in reduced free volume elements and cavity sizes. Reduced free volume in crosslinked PIM led to an increase in gas selectivity and decrease in gas permeability due to increase in diffusivity selectivity from low diffusion coefficients [192]. All crosslinking methods improve structural stability and resistance to CO2 swelling-induced densification. 3.3.3. PIMs with substituted pendant groups Functionalized PIMs have been investigated by introducing various pendent groups, basically on nitrile ( CN) sites [118,194–196]. The pendent phenylsulfone groups were substituted via a post-polymerization method to increase the selectivity while sacrificing gas permeability. Sulfone groups are known to raise the Tg and reduce free volume elements by space filling. Eventually, gas permeability is also reduced. However, sulfone groups give an improvement in gas selectivity without losing thermal stability. PIM membranes containing trifluoromethyl and

phenylsulfone side groups (TFMPSPIM), as a series of ladder polymers, have been prepared by replacing the nitile sites of PIM-1 with trifluoromethyl ( CF3 ) group and phenylsulfone ( SO2 C6 H5 ) groups [194]. The gas transport performances of TFMPSPIM membranes, as expected, demonstrated an improvement in gas selectivity for O2 /N2 and CO2 /N2 [194]. Other modifications of PIM series have been followed by DSPIM series where the nitrile groups in PIM-1 were substituted by various sulfone groups. Such modifications were expected to yield high gas selectivity with chain rigidity. DSPIM membranes also exhibited high gas selectivity for O2 /N2 and CO2 /N2 with reasonable gas permeability [195]. PIMs with substituted pendant groups modified by post-polymerization have been introduced as shown in Fig. 20 [118,195–198]. PIMs with CO2 -philic pendant tetrazole groups (TZPIMs) have been investigated by cycloaddition reaction (click reaction) between an inorganic azide and nitrile groups on PIM-1 [118,196]. The interaction between CO2 and N-containing organic heterocyclic molecules of tetrazole groups presented a string interaction, including hydrogen bonding with the oxygen atoms of CO2 . Consequently, the CO2 -philic parts present in PIMs were investigated. TZPIM showed different solution properties from PIM-1, where TZPIM became insoluble in CH2 Cl2 , THF, and CHCl3 . Meanwhile, the modified PIMs

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Fig. 20. PIMs with substituted pendant groups by post-polymerization modification [118,195,197,198].

were soluble in polar aprotic solvents such as DMF, DMAc, and NMP. Therefore, TZPIMs are also solution-processable from those solvents to prepare membrane films. CO2 -philic functional groups in PIMs could demonstrate a potential for both high CO2 permeability, as well as high selectivity for CO2 /light gases [118]. TZPIM membranes exhibit outstanding gas separation performance with high gas permeability, especially a high CO2 permeability with high gas selectivity of CO2 /CH4 and CO2 /N2 . Moreover, their anti-plasticization properties during CO2 permeation are also substantial at high pressure conditions. The CO2 permeability of TZPIM membranes has demonstrated over 3000 Barrers with a CO2 /N2 selectivity of around 30, as shown in Table 2 [118]. Carboxylated PIM-1 and thioamide-group introduced PIM1 have also been investigated by simple hydrolysis of PIM-1 [195,197,198]. Carboxylated PIM-1 maintained good processability with enhanced solubility in common organic solvent. However, carboxylation of PIM-1 increased gas selectivity but decreased permeability due to the closed chain distance by hydrogen or intermolecular bonding in carboxylated PIMs [195,197]. Thioamide-PIM-1 has been introduced by phosphorus pentasulfide as a thionating agent. The change in tendency of gas transport property after modification was similar to that of carboxylated PIM1. Gas permeability of thioamide-PIM-1 was reduced while its gas selectivity was increased after modification of PIM-1. Ethanol soaking improved gas permeability by enhancing polymer free volume [198].

3.3.4. PIM blends and mixed matrix membranes (MMMs) Various modifications of PIM membranes have been investigated in a way of forming polymer blend with other polymers or inorganic materials [199–205]. Commercially available Matrimid was selected to prepare polymer blend with PIM and PIM-1/Matrimid blends exhibit partially miscible behavior. Excellent thermal stability with high Tg of Matrimid increased Tg of PIM-1/Matrimid blends and gas selective property of Matrimid increased gas selectivity of PIM-1/Matrimid blends for O2 /N2 and CO2 /CH4 separation. PIM-1/Matrimid (70:30) blend membranes demonstrated CO2 permeability of 558 Barrers with CO2 /CH4 selectivity of 25 [201]. Moreover, these blends with high contents of Matrimid (15:85) have been fabricated to hollow fiber membranes with very thin and defect-free effective layer and exhibited CO2 permeance of 243 GPU with CO2 /CH4 selectivity of 34 [199]. Diamine cross-linking has also been introduced to modify PIM-1/Matrimid blends with high H2 /CO2 selectivity. Various diamine monomers such as Ethylenediamine (EDA), Triethylenetetramine (TETA), and 1,4-diaminobutane (BuDA) were used and cross-linking time was also controlled. PIM-1/Matrimid (90:10) blend membranes cross-linked with TETA for 2 h demonstrated H2 permeability of 395 Barrers with a H2 /CO2 selectivity of 9.6 at 35 ◦ C and 3.5 atm [200]. Polymer blends with inorganic materials, so-called mixed matrix membranes (MMM), have been investigated with various materials such as silica nanoparticle, zeolite

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Fig. 21. Benzodioxane-containing dianhydride and bulky diamines for PIM-PIs [182,210].

imidazolate frameworks (ZIFs), carbon nanotubes (CNTs), porous organic cage molecules, and silicalite-1 [202–207]. Mixed matric membranes are usually prepared to improve their gas permeation or separation properties of polymeric membranes. Fumed silica nanoparticles in PIM-1 increased gas transport performances because silica nanoparticles created void cavities as a permeable space in membrane matrix. However, it decreased gas selectivity. PIM-1 of fumed silica-filled nanocomposites with silica loadings of 19.1% demonstrated CO2 permeability of 12,182 Barrers with CO2 /N2 selectivity of 7.8 [203]. Single-walled and multi-walled carbon nanotubes (SWCNTs and MWCNTs)

also improved gas permeability of polymer membranes. CNTs that were functionalized with polyethylene glycol (PEG) were dispersed into PIM-1 matrix loading from 0.5 to 3.0 wt%. These composite membranes exhibited significant improvement in gas permeability and also relevant gas selectivity [204]. PIM-1 membranes with ZIF-8 as inorganic filler demonstrated an increase in gas permeability from well-defined pore size of 3.4 A˚ in ZIF-8. Controlling the loading of ZIF-8, composite membranes demonstrated high gas permeability with good selectivity. PIM-1 membrane with ZIF-8 28% demonstrated CO2 permeability of 4270 Barrers with CO2 /CH4 selectivity of 18.6 [205]. Porous

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Fig. 22. Preparation of PIM-PI-OH and PIM-TR-PBO [209,212].

crystalline solids were used as porous filler in mixed matrix membranes which provided good transport properties. Introduction of porous imine cage CC3 to PIM-1 enhanced gas permeability. PIM-1 MMM membranes with 30% CC3 exhibited CO2 permeability of 5430 Barrers where permeability gradually increased with the amount of CC3 loading from 10% to 30%. In this case, ethanol soaking also improved gas permeability. CO2 permeability of PIM-1 MMM membranes with 30% CC3 after ethanol treatment reached 37,400 Barrers. This approach can also be introduced for other porous cage organic molecules including calixarenes or cucurbiturils [207]. Silicalite-1 has been widely studied as microporous materials such as zeolite and silica nanoparticle. Mixed matrix membranes usually showed high gas permeability because filler acted as a gas permeation path. However, an introduction of silicalite-1 crystals with smaller free volume than PIM-1 decreased the gas permeability. CO2 permeability of mixed matrix membrane with silicalite-1 was reduced from 4390 Barrers to 2530 Barrers, but gas selectivity was improved [206]. 3.3.5. Polyimides of intrinsic microporosity (PI-PIMs) Polyimides with a ladder-like and contorted structure, which is observed in PIM series, have been introduced to improve the physical properties of PIMs [182,208–211]. Those polyimides, called PIM-PI, shown in Fig. 21, were prepared from benzodioxane-containing dianhydride and bulky diamines to acquire both the microporosity of PIMs and advanced physical properties of polyimides [182,208,210]. PIM-PI membranes demonstrated outstanding gas permeability compared with other conventional polyimide membranes. They also exhibited good chemical and thermal stability from polyimide, as well as good solubility for most polar aprotic solvents, which can increase solution-processibility for membrane formation. PIM-PIs possess the advantages of two polymers (i.e., PIM and PI). One example of PIM-PI membranes showed 490 Barrers of O2 permeability with an O2 /N2 selectivity of 3.5 [182]. Another approach to prepare PIM-PI was performed using hydroxyl-functionalized monomers in the polyimide part [209]. Hydroxyl groups in polymer

backbones could significantly increase the CO2 solubility, which will improve CO2 permeability. In addition, their hydrophilicity caused by hydroxyl groups could present a high water vapor sorption capacity, which will be beneficial in gas dehydration applications. The CO2 permeability of PIM-PI-OH demonstrated 263 Barrers with a CO2 /CH4 selectivity of 29. PIM-PI-OH membranes also show better plasticization resistance at a high CO2 partial pressure [209]. Although the permeability of copolymer membranes is one order of magnitude less than original PIM membranes, permeance can be improved by reducing the effective thickness of the asymmetric membrane or hollow fibers. In commercial scale hollow fiber membranes for gas separation application, the usual effective thickness of the skin layer is about 0.04 ␮m, meaning that 300 Barrers of gas permeability can achieve 7500 GPU for membrane module permeance. 3.3.6. TR polymers with spirobisindane (PIM-TR-PBO) A recent study on structural modification has investigated combining the structures of TR polymers and PIMs, e.g., producing a so-called PIM-TR-PBO or spiroTR-PBO [212,213]. Diamine monomers containing hydroxyl functional groups and spiro centers have been introduced to prepare PIM-TR-PBO. At the beginning, PIM-PI-OH was prepared, which had a similar or identical structure of a previous study [209]. Then, thermal rearrangement was performed to obtain the PIM-TR-PBO structure shown in Fig. 22. Hydroxyl groups in PIM-PI-OH aided the thermal conversion from polyimide into polybenzoxazole with a ring closing thermal reaction. Free volume elements of PIM-TR-PBO increased from the precursor PIM-PI-OH by 3–30% depending on the chemical structure, and thus their gas permeability was also improved. The CO2 permeability of PIM-TR-PBO membranes demonstrated 675 Barrers with a CO2 /CH4 selectivity of 20 [212]. Moreover, their mixed-gas CO2 /CH4 separation property was measured and demonstrated 557 Barrers of CO2 permeability with selectivity of 15 [213]. The advantage of the structural combination of TR polymer and PIMs is centered on their mechanically robust property. The average tensile strength

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Fig. 23. (a) Structure of kink linkage of TR polymers. (b) Angle distribution of kink linkage in Spiro segment (red) and 6F segment (black). (c) Relationship of elongation at break and fractional free volume of various glassy polymers, spiroTR-PBOs (䊉red) and spiroHPIs (䊉blue) [212]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.) Copyright 2013. Reproduced with permission from Elsevier Ltd.

Fig. 24. Upper bound relationship for O2 /N2 , CO2 /N2 , and CO2 /CH4 separation: (a) TR-␣-PBO, (b) TR-␣-PEBO, (C) TR-␣-PBI, (d) TR-␤-PBO, (e) TR-␣-PBOco-PPL, (f) TR-␣-PBO-co-PI, (g) TR-␣,␤-PBO, (h) thermally labile PI, (i) basic PIMs, (j) rigid PIMs, (k) cross-linked PIMs, (l) PIMs with substituted pendant groups, (m) PI-PIM, (n) PIM-TR-PBO, and (o) PIM blends.

and elongation at break of PIM-TR-PBO membrane achieves over 80 MPa and 20%, respectively, which is about the highest mechanical strength among polymers with high free volume elements, as shown in Fig. 23. The degree of entanglement of molecular chains by the kink in spiro segment could increase elongation at break. The robust mechanical property of PIMs provides easy processability for potential scale-up in gas separation and other applications [212]. Various modifications of PIMs have been performed to investigate the resulting gas transport properties. Table 2 summarizes gas permeability and selectivity of PIM-based polymer membranes. Gas transport properties of PIM membranes varied even though they had the same chemical structure because most of reported PIMs were synthesized under different conditions. Moreover, methanol treatment after membrane formation swelled polymer matrix and improved initial gas permeability of the membrane. Gas transport properties of PIM membranes depended on chain rigidity and stiffness of the resulting polymers. 4. Future outlooks and conclusions Recent development in high performance polymer membranes has advanced the gas transport performance

of the first generation gas separation membranes. Polymer rigidity is one of the key issues to improve membrane gas transport performances, which stems from an intrinsic microporosity with high free volume elements. Microporous polymers have been of interest as gas separation membrane materials, as well as gas storage materials, due to their high porosity and large surface area. Various classes of microporous polymers, including PCPs, HCPs, CMPs, and COFs, have been developed for their applications in gas storage materials. However, most of them have insoluble network due to rigid and cross-linked network and it limits their application as membrane materials. Based on the high porosity with rigid structures, they have a very high potential for application in gas separation membranes. Solubility of microporous organic materials can be improved by diminution of cross-linked or branched network and improved processability allows them to be applied for gas separation membrane process. Recently developed TR polymer and PIM membranes have received much attention due to their extraordinary gas permeability for light gases. Their gas separation performances have surpassed or are close to the current performance limits of polymeric membranes, the so called the upper-bound as shown in Fig. 24. This new series of microporous polymer membranes is expected to be utilized in large scale gas

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separation applications by module fabrication using their solution-processability. They are also expected to meet the economic cost requirement of the gas separation process, as well as the separation efficiency, with their outstanding gas permeability. Future developments of microporous membranes for gas separation will require enhanced sorption of target gas pairs while maintaining the extraordinary gas permeability due to their high free volume elements. Tuning the cavity sizes is quire essential for target gas pairs in order to increase membrane performance. Cavity sizes larger than 4 A˚ will have a great potential for the separation of small hydrocarbons such as C2 , C3 , and C4+ . The petrochemical industry is waiting for high-performance membranes to separate mixtures of these hydrocarbons. Future direction will focus on the separation of these hydrocarbons by understanding the capability in this field.

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This research was supported by the Korea CCS R&D Center (KCRC) (2014M1A8A1049305), and the Nano Materials Technology Development Program (2012M3A7B4049745), through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning. References [1] Xie Y, Wang TT, Liu XH, Zou K, Deng WQ. Capture and conversion of CO2 at ambient conditions by a conjugated microporous polymer. Nat Commun 2013;4:1960–6. [2] Vilela F, Zhang K, Antonietti M. Conjugated porous polymers for energy applications. Energy Environ Sci 2012;5:7819–32. [3] Cooper AI. Conjugated microporous polymers. Adv Mater 2009;21:1291–5. [4] Chen L, Honsho Y, Seki S, Jiang D. Light-harvesting conjugated microporous polymers: rapid and highly efficient flow of light energy with a porous polyphenylene framework as antenna. J Am Chem Soc 2010;132:6742–8. [5] Jiang JX, Su F, Trewin A, Wood CD, Campbell NL, Niu H, Dickinson C, Ganin AY, Rosseinsky MJ, Khimyak YZ, Cooper AI. Conjugated microporous poly(aryleneethynylene) networks. Angew Chem Int Ed 2007;46:8574–8. [6] Dillon AC. Carbon nanotubes for photoconversion and electrical energy storage. Chem Rev 2010;110:6856–72. [7] Wu D, Xu F, Sun B, Fu R, He H, Matyjaszewski K. Design and preparation of porous polymers. Chem Rev 2012;112:3959–4015. [8] Yang Z, Zhang J, Kintner-Meyer MCW, Lu X, Choi D, Lemmon JP, Liu J. Electrochemical energy storage for green grid. Chem Rev 2011;111:3577–613. [9] Graetz J. New approaches to hydrogen storage. Chem Soc Rev 2009;38:73–82. [10] McKeown NB, Budd PM. Polymers of intrinsic microporosity (PIMs): organic materials for membrane separations, heterogeneous catalysis and hydrogen storage. Chem Soc Rev 2006;35:675–83. [11] Bu XH. Microporous organic polymers for gas storage and separation applications. Phys Chem Chem Phys 2013;15:5430–42. [12] Couture G, Alaaeddine A, Boschet F, Ameduri B. Polymeric materials as anion-exchange membranes for alkaline fuel cells. Prog Polym Sci 2011;36:1521–57. [13] Li Q, Jensen JO, Savinell RF, Bjerrum NJ. High temperature proton exchange membranes based on polybenzimidazoles for fuel cells. Prog Polym Sci 2009;34:449–77. [14] Park CH, Lee CH, Guiver MD, Lee YM. Sulfonated hydrocarbon membranes for medium-temperature and low-humidity proton exchange membrane fuel cells (PEMFCs). Prog Polym Sci 2011;36:1443–98. [15] Xiao Y, Low BT, Hosseini SS, Chung TS, Paul DR. The strategies of molecular architecture and modification of polyimide-based

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