Polymer–inorganic nanocomposite membranes for gas separation

Separation and Purification Technology 55 (2007) 281–291

Review

Polymer–inorganic nanocomposite membranes for gas separation Hailin Cong, Maciej Radosz, Brian Francis Towler, Youqing Shen ∗ Soft Materials Laboratory, Department of Chemical & Petroleum Engineering, University of Wyoming, Laramie, WY 82071-3295, USA Received 1 September 2006; received in revised form 20 December 2006; accepted 20 December 2006

Abstract Polymer–inorganic nanocomposite membranes present an interesting approach to improve the separation properties of polymer membranes because they possess properties of both organic and inorganic membranes such as good permeability, selectivity, mechanical strength, and thermal and chemical stability. The preparations and structures of polymer–inorganic nanocomposite membranes, their applicability to gas separation and separation mechanism are reviewed. © 2007 Elsevier B.V. All rights reserved. Keywords: Nanocomposite membranes; Polymer membrane; Gas separation; Gas transport mechanism

Contents 1. 2. 3.

4. 5.

6.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Types of nanocomposite membrane by structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Preparation of nanocomposite membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Solution blending . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. In situ polymerization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3. Sol–gel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Gas separation properties of nanocomposite membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Gas transport mechanisms in nanocomposite membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. Maxwell’s model . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. Free-volume increase mechanism . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3. Solubility increase mechanism . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.4. Nanogap hypothesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions and future directions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction In the last two decades significant improvements in the performance of polymeric membranes for gas separation have been made [1–6], and our understanding of the relationships among the structure, permeability and selectivity of polymeric membranes has been greatly advanced [2,3]. Newer polymeric membrane materials such as polyimides (PI) and cross-linked

∗

Corresponding author. Tel.: +1 307 7662468; fax: +1 307 7666777. E-mail address: [email protected] (Y. Shen).

1383-5866/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.seppur.2006.12.017

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polyethylene glycol (PEG) have been continuously developed [7–11]. Some polymeric membranes have already been used in industry [12,13]. For instance, a plant separating air into its constituent gases and producing pure nitrogen at nearly 24 t h−1 in Belgium by Praxair Co. began operation in 1996 [3]. Polymeric membranes tend to be more economical than other membranes because of their ability to be easily spun into hollow fibers or spiral-wound modules due to their flexibility and solution processability [1]. Despite these advantages and progresses, polymeric membranes are still restricted by the trade-off trend between gas permeability and selectivity, as suggested by Robeson [14].

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Nomenclature A preexponential factor AF2400 poly(2,2-bis(trifluoromethyl)-4,5-difluoro-1,3dioxole-co-tetrafluoroethylene) AIBN 2,2 -azobisisobutyronitrile APrMDEOS aminopropylmethyldiethoxysilane APrTMOS aminopropyltrimethoxysilane BPPOdm poly(2,6-dimethyl-1,4-phenylene oxide) BPPOdp poly(2,6-diphenyl-1,4-phenylene oxide) C60 fullerene CNTs carbon nanotubes COOH-CNTs carboxylic acid-functionalized CNTs D diffusivity or diffusion coefficients DA diffusivity of gas A DABA diaminobenzoic acid DSC differential scanning calorimetry Ed energy of diffusion Ep apparent activation energy 6FDA hexafluoroisopropylidene diphthalic anhydride FPAI fluorinated poly(amide-imide) 6FpDA hexafluoroisopropylidene dianiline Hs enthalpy of sorption MA methacrylic acid MTMOS methyltrimethoxysilane NMP N-methylpyrrolidone P permeability PA permeability of gas A gas permeability in composite membrane Pc Pp gas permeability in pure polymer matrix P0 preexponential factor PA polyamide PAA polyamic acid PAI poly(amide-imide) PALS position annihilation lifetime spectroscopy PAN polyacrylonitrile PEBAX poly(amide-6-b-ethylene oxide) PEG polyethylene glycol PEI poly(ether imide) PEO poly(ethylene oxide) PI polyimides PMA poly(methacrylic acid) PMDA pyromellitic dianhydride PMP poly(4-methyl-2-pentyne) PPEPG PPG-block-PEG-block-PPG PPG poly(propylene glycol) PSF polysulfone PTMOS phenyltrimethoxysilane PTMSP poly(1-trimethylsilyl-1-propyne) R ideal gas constant S solubility SA solubility of gas A SEM scanning electron microscopy T absolute temperature Tg glass transition temperature TEM transmission electron microscopy

TEOS TMOS Vf V* WAXD XRD

tetraethoxysilane tetramethoxysilane average free volume minimum free volume element size wide-angle X-ray diffraction X-ray diffraction

Greek symbols αA/B permeability selectivity of gas A to B γ overlap factor Φf volume fraction of nanofiller in the membrane Modifications of the chemical structure of a polymer often lead to an improvement in permeability at the cost of selectivity, or vice versa [3]. Additionally, the segmental flexibility of polymeric membranes often limits their ability to discriminate similar-sized penetrants and they often lose performance stability at high temperatures [1]. On the other hand, inorganic membrane materials such as molecular sieving materials usually rely on a difference in molecular size to achieve separation. On a laboratory scale, these membranes show extremely attractive gas permeation and separation performance [15,16]. However, it is still difficult and expensive to fabricate large membranes due to their fragile structures [1,2,17]. Therefore, polymeric membranes are still attractive, but alternate approaches that can enhance their gas separation properties well above the Robeson line are needed. Polymer–inorganic nanocomposite materials, herein defined as inorganic nanofillers dispersed at a nanometer level in a polymer matrix, have been investigated for gas separation, and have the potential to provide a solution to the trade-off problem of polymeric membranes [18,19]. For example, many polymer–inorganic nanocomposite membranes show much higher gas permeabilities but similar or even improved gas selectivities compared to the corresponding pure polymer membranes [20–26]. The nanocomposite materials may combine the advantages of each material: for instance, the flexibility and processability of polymers, and the selectivity and thermal stability of the inorganic fillers. Additionally, the gas separation performance of nanocomposite membranes can be further enhanced by chemical modification [27]. For instance, the introduction of organic functional groups on an inorganic filler surface sometimes contributes to not only a better dispersion of the inorganic material in the polymer membrane, but also a better absorption and transportation of penetrants, which results in favorable selectivity and permeability [27,28]. Membrane structure can be controlled by either the degree of cross-linking of the polymer matrix, or the types of connection bonds between the polymer and inorganic phases in the nanocomposite material [28,29]. 2. Types of nanocomposite membrane by structure As shown in Fig. 1, polymer–inorganic nanocomposite membranes can be divided into two types according to their structure: (a) polymer and inorganic phases connected by covalent bonds

H. Cong et al. / Separation and Purification Technology 55 (2007) 281–291

Fig. 1. Illustration of different types of polymer–inorganic nanocomposite membranes. (a) Polymer and inorganic phases connected by covalent bonds and (b) polymer and inorganic phases connected by van der Waals force or hydrogen bonds.

and (b) polymer and inorganic phases connected by van der Waals force or hydrogen bonds [30]. 3. Preparation of nanocomposite membranes Because of the huge difference between the polymer and inorganic materials in their properties and strong aggregation of the nanofillers, polymer–inorganic nanocomposite membranes cannot be prepared by common methods such as melt blending and roller mixing. The most commonly used preparation technologies for the fabrication of nanocomposite membranes can be divided into the following three types [31]. 3.1. Solution blending Solution blending is a simple way to fabricate polymer– inorganic nanocomposite membranes. A polymer is first dissolved in a solvent to form a solution, and then inorganic nanoparticles are added into the solution and dispersed by stirring. The nanocomposite membrane is cast by removing the solvent. For example, Genne et al. [32] prepared polysulfone (PSF)/ZrO2 nanocomposite membranes using 18 wt.% PSF solution in N-methylpyrrolidone (NMP) with adding various amounts of ZrO2 nanoparticles. The membrane permeability increased as the ZrO2 weight fraction increased. Wara et al. [33] reported the fabrication of nanocomposite membranes of cellulose/Al2 O3 by using the solution blending. The solution blending method is easy to operate and suitable for all kinds of inorganic materials, and the concentrations of the polymer and inorganic components are easy to control [34]. However, the inorganic ingredients are liable to aggregate in the membranes [35,36]. 3.2. In situ polymerization In this method, the nanoparticles are mixed well with organic monomers, and then the monomers are polymerized. There are

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often some functional groups such as hydroxyl, carboxyl on the surface of inorganic particles, which can generate initiating radicals, cations or anions under high-energy radiation, plasma or other circumstances to initiate the polymerization of the monomers on their surface. For instance, nanocomposite membranes of poly(methacrylic acid) (PMA)/TiO2 were synthesized from TiO2 nanopowder/methacrylic acid dispersions under microwave radiation [37]. Doucoure et al. [38] reported in situ plasma polymerization of fluorinated monomers on mesoporous silica membranes. Patel et al. [39,40] prepared cross-linked nanocomposite membranes of PEG/silica and poly(propylene glycol) (PPG)/silica by dispersing silica nanoparticles in diacrylate-terminated PEG and PPG, and subsequent radical polymerization initiated by 2,2 -azobisisobutyronitrile (AIBN). Nunes et al. [41] reported the fabrication of nanocomposite membranes of poly(ether imide) (PEI)/SiO2 by using in situ polymerization. In the in situ polymerization method, inorganic nanoparticles with functional groups can be connected with polymer chains by covalent bonds. However, it is still difficult to avoid the aggregation of inorganic nanoparticles in the formed membranes. 3.3. Sol–gel The sol–gel method is the most widely used preparation technology for nanocomposite membranes. In this method, organic monomers, oligomers or polymers and inorganic nanoparticle precursors are mixed together in the solution. The inorganic precursors then hydrolyze and condense into well-dispersed nanoparticles in the polymer matrix. The advantage of this method is obvious: the reaction conditions are moderate—usually room temperature and ambient pressure, and the concentrations of organic and inorganic components are easy to control in the solution. Additionally, the organic and inorganic ingredients are dispersed at the molecular or nanometer level in the membranes, and thus the membranes are homogeneous [42–45]. For example, Iwata et al. [46] reported that by using the sol–gel method, a nanocomposite membrane of polyacrylonitrile (PAN) with hydrolysate of tetraethoxysilane (TEOS) as the inorganic phase showed a good performance in O2 /N2 separation. Gomes et al. [47] prepared nanocomposite membranes of poly(1-trimethylsilyl-1-propyne) (PTMSP)/silica by sol–gel copolymerization of TEOS with different organoalkoxysilanes in the tetrahydrofuran solution of PTMSP. Fig. 2 shows the scanning electron microscopy (SEM) images of cross-sections of PTMSP membrane (a) and PTMSP/silica membrane (b) prepared by sol–gel process. 4. Gas separation properties of nanocomposite membranes The permeability (P) of a gas through a membrane is proportional to the solubility (S) and diffusivity (D) of the gas in the membrane (P = D × S). Thus, adding inorganic nanofillers may affect the gas separation in two ways: the interaction between polymer-chain segments and nanofillers may disrupt

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Fig. 2. SEM photomicrographs of cross-sections of: (a) PTMSP membrane and (b) PTMSP/silica nanocomposite membrane prepared by sol–gel process [47]. (Reproduced with permission from Elsevier Co.)

the polymer-chain packing and increase the voids (free volumes) between the polymer chains, and thus enhance gas diffusion [20,22,23,48–50]; the hydroxyl and other functional groups on the surface of the inorganic phase may interact with polar gases such as CO2 and SO2 , improving the penetrants’ solubility in the nanocomposite membranes [26,51]. Various combinations of polymers and nanofillers have been tested for gas separations. The results are summarized in Table 1. Among these nanocomposite membranes, polyimide/silica materials have received the most attention for the gas permeation studies [28,52,53]. Joly et al. [21,54] fabricated the polyimide/silica membranes containing 32 wt.% silica via the sol–gel method by adding tetramethoxysilane (TMOS) to polyamic acid (PAA) solution and subsequently imidizing at 60–300 ◦ C. As shown in Table 1, the nanocomposite membrane had a higher permeability for CO2 (PCO2 = 2.8 Barrer) and a greater CO2 /N2 selectivity (αCO2 /N2 = 22) compared to the polyimide membrane (PCO2 = 1.8 Barrer; αCO2 /N2 = 18). The gas permeation results were analyzed using the dual sorption model. In this model, it was assumed that the gas molecules dissolved in the polymer could be classified into two distinct populations: (a) Henry-type dissolution and (b) Langmuir-type sorption. The authors attributed the increased gas permeation of the nanocomposite membrane compared to the polyimide membrane to enhanced gas solubility due to an increased contribution of Henry’s type dissolution. Using X-ray diffraction (XRD) and SEM, the authors showed that the addition of TMOS to the PAA induced some morphological modifications in the polymer matrix. Kusakabe et al. [55] reported that the CO2 permeability in a polyimide/SiO2 hybrid nanocomposite membrane was 15 times lager than that in the corresponding polyimide. The permselectivity of CO2 to N2 was 25 at 30 ◦ C. Contributions of the silica and polyimide phases to the composite membrane’s permeance

were analyzed using a two-phase permeation model. The effective thickness of the rate-controlling polyimide phase was less than one tenth of the thickness of the composite membrane. Homogeneous nanocomposite membranes of polyimide– siloxane copolymers containing different silica contents were prepared by Smaihi et al. [56] via the sol–gel process of pyromellitic dianhydride (PMDA), aminoalkoxysilane, and TMOS. They used two coupling agents, aminopropyltrimethoxysilane (APrTMOS) and aminopropylmethyldiethoxysilane (APrMDEOS) to provide bonding between the imide and the inorganic silica. Higher gas permeability was observed for the membrane using APrMDEOS than for the membrane using APrTMOS at the same silica content. IR studies of the nanocomposite material revealed that the presence of methyl side groups linked to the silicon of the APrMDEOS precursor inhibited the formation of −OH linked bonds in the material. Moaddeb and Koros [22] studied the gas transportation properties of thin polyimide membranes in the presence of silica particles. In the nanocomposite membranes on silicaimpregnated aluminum oxide substrates, the presence of silica improved the gas separation properties of the polyimide layer, particularly for O2 and N2 . The increase in permeability was due to silica-disrupting the polymer-chain packing. The observed significant increase in the glass-transition temperature (Tg ) suggested the restriction of chain segmental mobility possibly due to adsorption of the polymer to the silica surface. Cornelius and co-workers [23,57–59] studied the effects of alkoxysilanes, their loading, and the morphology of the resulting polyimide/silica nanocomposite membranes on the permselectivity of several gases including CO2 , N2 , and CH4 . The polyimides were prepared from hexafluoroisopropylidene diphthalic anhydride (6FDA)–hexafluoroisopropylidene dianiline (6FpDA)–diaminobenzoic acid (DABA) and the alkoxysilanes included phenyltrimethoxysilane (PTMOS),

[54] [56] [23] [24] [25] [53] [51] [69] [83] [65] [26] [39] [70] [71] [48] [46] [47] [63] 4.3 14.9

3.2

188

9.1

0.65 21 29 29 79

2.4 2.6

Permeability of CO2 , 1 Barrer = 10−10 cm3 (STP) cm/cm2 s cmHg. Permselectivity of CO2 to N2 . a

b

4234 1157

7.36

1.3 34.9 27.3

66.3

11.3

2.46 1.4 0.56 0.17

24.9 15.0 5.34 3.52

50 15.04

0.76 2.61

15 42.9 523 436 155 277 67 94.2 125

2.16 0.08

10.3 0.2

Silica Silica Silica Silica TiO2 Silica TiO2 Silica Silica CNT Silica Silica Silica Silica C60 Silica Silica Silica PI PI PI PI PI PI PAI BPPOdm BPPOdp BPPOdp PEBAX PEG PEG PPEPG PS PAN PTMSP PMP

2.8 41 80 19

38.3 89

33 15 16

37 238

4.0 14

22 5.3 16 41

0.13 7.74 5 0.46 0.08 0.3

PO2 (Barrer) PH2 (Barrer) PCH4 (Barrer) PCO2 a (Barrer) Nanofiller Polymer matrix

Table 1 Gas-separation performance data for representative polymer–inorganic nanocomposite membranes

PN2 (Barrer)

αCO2 /N2 b

αCO2 /CH4

αCO2 /H2

24.0 22.3

αC4 H10 /CH4

9.5 8.7

αO2 /N2

αH2 /N2

1.33

51

αH2 /CH4

Refs.

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methyltrimethoxysilane (MTMOS), TMOS, and TEOS. These nanocomposite membranes were annealed at 400 ◦ C to drive the sol–gel reactions to a greater extent. In general, the annealing process increased the gas permeation of the nanocomposite membranes by about 200–500%, while the permselectivity dropped slightly. An exception was the 6FDA–6FpDA–DABA25 (containing 25 mol% DABA) membrane. The nanocomposite membranes with 22.5 wt.% TMOS and MTMOS both had increased CO2 permeabilities (see Table 1) and CO2 /CH4 permselectivity [23]. The authors attributed the increase in gas permeation to changes in the free volume distribution and enhanced local segment mobility of the chain ends resulted from the removal of sol–gel condensation and polymer degradation by-products. Suzuki and Yamada [24] reported the physical and gas transport properties of a 6FDA-based hyperbranched polyimide/silica nanocomposite membrane prepared using polyamic acid, water and TMOS via the sol–gel technique. CO2 , O2 and N2 permeability coefficients of the membrane increased as the silica content increased. It was pointed out that the increased gas permeabilities were mainly attributable to the increase in the gas solubilities. In contrast, CH4 permeability of the nanocomposite membranes decreased with increasing silica content because of the decrease in the CH4 diffusivity. As a result, CO2 /CH4 selectivity (see Table 1) of the nanocomposite membranes increased remarkably. This kind of nanocomposite membrane had high thermal stability and excellent gas selectivity, and is expected to be a high-performance gas-separation membrane. Polymer–inorganic nanocomposite membranes based on a fluorinated poly(amide-imide) (FPAI) and TiO2 were fabricated by Hu et al. [51] via the sol–gel method. An aromatic poly(amide-imide) (PAI) was chosen as the polymer matrix material because it provided superior mechanical properties, high thermal stability, solvent resistance, and high permeability. The nanocomposite membrane had a more rigid and dense structure than the corresponding pure FPAI membrane. The authors observed a specific interaction between such gases as CO2 and H2 and the TiO2 particles in the nanocomposite membrane. Higher selectivities for CO2 /CH4 and H2 /CH4 gas pairs (see Table 1) were observed in the composite membrane containing a low concentration (7.3 wt.%) of TiO2 . Kong et al. [25] prepared polyimide/TiO2 nanocomposite membranes by blending TiO2 sol and a polyimide solution. Because of the improved TiO2 -sol preparation process and blending method, the TiO2 content could reach about 40 wt.% in the membrane. There existed a strong interaction between the TiO2 phase and the polyimide phase. A higher TiO2 content in the membrane resulted in a greater enhancement of the gasseparation performance. The H2 and O2 permeabilities of the membrane with 25 wt.% TiO2 were 14.1 and 0.72 Barrer, respectively, which were 3.7 times and 4.3 times higher than those of the pure polyimide. The selectivities of H2 /N2 and O2 /N2 (see Table 1) were also slightly improved compared to the pure polymer membrane (αH2 /N2 = 167, αO2 /N2 = 9.3). Polymer–inorganic nanocomposite membranes of poly (amide-6-b-ethylene oxide) (PEBAX)/silica were prepared by Kim and Lee [26] via in situ polymerization of TEOS using

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the sol–gel process. The membrane containing 27 wt.% silica had a CO2 permeability of 277 Barrer and CO2 /N2 selectivity of 79, both higher than those of the pure polymer membrane. The PEBAX copolymer consisted of two distinct regions, the impermeable crystalline polyamide (PA) phase, and the permeable amorphous poly(ethylene oxide) (PEO) phase [60]. Using wide-angle X-ray diffraction (WAXD) and differential scanning calorimetry (DSC), the authors showed that the presence of silica domains in the nanocomposite membrane significantly decreased the degree of crystallinity of the PA phase and caused reorientation of the PEO phase. They reported that the nanocomposite membranes exhibited higher gas permeability coefficients and permselectivities than the PEBAX alone, particularly at an elevated temperature. The gas permeability and permselectivity increased as the silica content increased in the nanocomposite membrane. It was concluded that the increases in permeability and permselectivity of the nanocomposite membrane arose from the strong interaction between CO2 molecules and the residual hydroxyl groups on the silica domain, and the additional sorption sites in polyamide block of PEBAX. Higuchi et al. [48] observed an increase in gas permeability upon adding fullerene (C60 ) particles to polystyrene. Permeability increased 41% for ethylene, 47% for N2 and 75% for ethane in a film containing 10 wt.% fullerene particles at 25 ◦ C. Selectivity decreased about 25% for O2 /N2 (see Table 1) and ethylene/ethane. The enhancement in permeability was attributed to increases in diffusion coefficients caused by the increased free volume in the membrane. Pinnau and He [61,62] and Merkel et al. [20,63] reported that adding nanosized impermeable particles of commercial fumed silica to poly(4-methyl-2-pentyne) (PMP) increased gas and vapor permeabilities with increased particle loading. For example, the n-butane permeability (see Table 1) increased by a factor of 3 relative to that of the pure PMP when 30 wt.% silica was added at 25 ◦ C. Silica particles did not alter the solubility of the nanocomposite, but they did significantly increase the gas diffusion coefficient. For instance, the CH4 diffusion coefficient doubled at 30 wt.% silica loading [63]. The silica particles (∼12 nm) used in these studies were small enough to disrupt polymer-chain packing in the polymers, which resulted in an increase in polymer fractional free volume [20]. The free volume increase was characterized using density and position annihila-

tion lifetime spectroscopy (PALS) measurements. By increasing fractional free volume, gas diffusion coefficients and, in turn, gas permeability increased [49,64]. Particle loading influenced the gas transport properties in these nanocomposite membranes. For example, N2 permeability tripled as silica loadings in PMP increased from 5 to 25 vol.% [62]. The enhancement in n-butane permeability coincided with a substantial enhancement in nbutane/CH4 mixed gas selectivity. The n-butane/CH4 selectivity (see Table 1) doubled relative to that of the pure PMP at a 30 wt.% silica concentration. The enhanced n-butane permeability and n-butane/CH4 selectivity were attributed to the increased free volume of PMP by the silica-particles [62]. Recently, our group fabricated nanocomposite membranes of brominated poly(2,6-diphenyl-1,4-phenylene oxide) (BPPOdp ) and carbon nanotubes (CNTs) via the solution-blending method [65]. CNTs were chosen as the inorganic filler material because they are very effective in reinforcing polymeric materials and thus may provide better mechanical properties [66–68] to the membrane. Fig. 3 shows the transmission electron microscopy (TEM) photomicrographs of cross-sections of BPPOdp /single-wall CNT nanocomposite membranes prepared by solution-blending method. The composite membranes had an increased CO2 permeability but a similar CO2 /N2 selectivity (see Table 1) compared to the corresponding pure BPPOdp membranes (PCO2 = 78 Barrer; αCO2 /N2 = 30). The CO2 permeability increased with increasing the carbon nanotube content and reached a maximum of 155 Barrer at 9 wt.% single-wall CNTs, or 148 Barrer at 5 wt.% multiwall CNTs. The CO2 /N2 separation performance was not sensitive to the CNT diameter and length. However, the carboxylic acid-functionalized CNTs (COOH-CNTs), which were more homogeneously dispersed in BPPOdp , neither increased the gas permeability nor deteriorated the gas separation performance. We concluded that due to the incompatibility of pristine CNTs and the BPPOdp chains, the polymer chains did not attach to the CNT wall tightly, forming narrow gaps surrounding the CNTs. Gas molecules thus easily passed through the gap and had a shortcut. This also explained why the addition of CNTs did not affect the CO2 /N2 selectivity. The modified CNTs surface was compatible with the polymer and the polymer chains could pack tightly on the CNT surface, closing the nanogap and leaving the gas permeability unenhanced. In another study, nanocompos-

Fig. 3. TEM photomicrographs of cross-sections of BPPOdp /single-wall CNT nanocomposite membranes prepared by solution-blending method.

H. Cong et al. / Separation and Purification Technology 55 (2007) 281–291

ite membranes of brominated poly(2,6-dimethyl-1,4-phenylene oxide) (BPPOdm ) and silicas were fabricated successfully via the solution-blending method [69], and we also found the composite membranes had an increased CO2 permeability but a similar CO2 /N2 selectivity (see Table 1) compared to the corresponding pure BPPOdm membranes. The permeabilities of all the gases increased with increasing silica concentration. For example, the PCO2 of the BPPOdm /10 nm silica membrane was 187 Barrer at 9 wt.% silica, and reached 523 Barrer at 23 wt.% silica, about five times of that of the pure BPPOdm membrane; in the same membranes the selectivity over N2 remained unchanged. The gas-separation performance of the present nanocomposite membranes can be further enhanced by modification of the fillers and matrices. For example, Patel et al. [27] studied the effects of nanoparticle functionality on CO2 -selective nanocomposite membranes derived from cross-linked PEG and found that methacrylate-functionalized silica nanoparticles were more effective in improving rheological properties and retaining high CO2 selectivity than the original hydroxyl-functionalized silica nanoparticles of comparable size in the cross-linked nanocomposite membranes. The reason for this difference was that the methacrylate-functionalized silica nanoparticles reacted with PEG-diacrylate oligomers in the cross-link process, and thus improved dispersion of fillers in the polymer matrix and increased the interaction between them, while the hydroxyl-functionalized silica nanoparticles could not attend the cross-link reaction. Kim et al. [70] reported that by using the sol–gel method, the nanocomposite membrane with PEG as the organic phase and hydrolysate of TEOS as the inorganic phase showed good performance in CO2 /N2 separation. The CO2 permeability was 94.2 Barrer with a CO2 /N2 selectivity of 38.3. By simply changing the PEG matrix to PPG-block-PEG-block-PPG (PPEPG), Sforca et al. [71] prepared nanocomposite membranes by the same method and reported that the CO2 permeability (see Table 1) increased to 125 Barrer, and the CO2 /N2 selectivity increased to 89. It was concluded that the introduction of PPG segments in the PEG chains not only disrupted the original polymer-chain packing but also changed the chemical affinities of penetrants in the matrix.

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The numerator represents the loss of membrane solubility due to the loss of polymer volume available for sorption. The denominator represents a decrease in diffusivity due to increasing the penetrant diffusion pathway length [63]. Both factors act to decrease permeability with increasing particle volume fraction [75]. Maxwell’s model partly explains the gas permeability loss in some nanocomposite membranes, especially in the fullerene-particle-filled polymer membranes [76–78]. However, in general, the addition of fullerene to polymers decreases permeability more than the loss predicted by Maxwell’s model [75]. Also, Higuchi et al. [48] observed an increase in gas permeability upon adding fullerene particles to polystyrene, whereas Maxwell’s model predicted a 14% loss in gas permeability for this system. Additionally, more and more polymer–inorganic nanocomposite membranes have been observed with similar non-Maxwellian effects [20,22,23,62,65]. For example, adding nanosized impermeable particles of commercial fumed silica (TS 530, 12 nm primary diameter) to a glassy polymer (e.g., PMP) increased gas and vapor permeabilities with increased particle loading (see Fig. 4) [62]. The problem with Maxwell’s model lies in its neglect of the interactions between the nanofillers and the polymer chains, and the nanofillers and the penetrants. In most nanocomposite membranes, these interactions are strong, and significantly change the diffusivity and solubility of penetrants. 5.2. Free-volume increase mechanism The effect of polymer free volume on penetrant diffusion coefficients is often modeled by the statistical-mechanical description of diffusion in a liquid of hard spheres proposed by Cohen and Turnbull [79]. This model provides the following expression for penetrant diffusion coefficients (D) [63]:   −γV ∗ (2) D = A exp Vf

5. Gas transport mechanisms in nanocomposite membranes 5.1. Maxwell’s model Adding impermeable inorganic nanoparticles to a polymer is typically expected to reduce the gas permeability [72]. Maxwell’s model, developed to analyze the steady-state dielectric properties of a diluted suspension of spheres [73], is often used to model permeability in membranes filled with roughly spherical impermeable particles [74]:   1 − Φf Pc = Pp (1) 1 + 0.5Φf where Pc and Pp are the permeability of the nanocomposite and the pure polymer matrix, respectively, and Φf is the volume fraction of the nanofiller.

Fig. 4. N2 permeability enhancement of PMP as a function of filler content at 25 ◦ C and 50 psig feed pressure [62]. (Reproduced with permission from Elsevier Co.)

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where A is a preexponential factor weakly dependent on temperature, γ an overlap factor introduced to avoid double-counting free volume elements, V* the minimum free volume element size that can accommodate a penetrant molecule (and is closely associated with penetrant size), and Vf is the average free volume in the media accessible to penetrants for transport. According to Eq. (2), an increase in polymer free volumes is expected to enhance penetrant diffusion. Based on the PALS measurements, Merkel et al. [49] reported that the addition of fumed silica increased the size of free-volume elements in poly(2,2-bis(trifluoromethyl)4,5-difluoro-1,3-dioxole-co-tetrafluoroethylene) (AF2400). The enhanced free volume of AF2400/silica resulted in augmented penetrant permeability and diffusion coefficients, similar to the observation in the PMP/silica nanocomposite membranes [20,63]. The authors concluded that the improvement in permeability reflected hybridization-induced disruption of polymer-chain packing and an accompanying elevated free volume available for molecular diffusion. Winberg et al. [50] studied the free volume in silica-filled PTMSP nanocomposite membranes with PALS at filler concentrations between 0 and 50 wt.%. A bimodal free-volume distribution was observed, and the size of larger free volume cavities was significantly increased upon addition of hydrophobic fumed silica. The authors observed a strong correlation between N2 permeability and the volume of the larger free-volume cavities in the nanocomposite membranes, and the permeability increased with increasing filler content. It is worth to mention that Hill [80] significantly extended the Cohen–Turnbull free volume theory recently by hypothesis that the accompanying increase in free volume in the nanocomposite membranes reflects a repulsive interaction between the polymer chains and inclusions during membrane casting. He proposed a theoretical model based on the Cohen–Turnbull statistical mechanical theory, which not only captured the correct dependence of the diffusive permeability and selectivity of polymeric nanocomposites on the inclusion size and volume fraction, but also achieved a quantitative interpretation of the Merkel’s experiments [20,63]. The free-volume increase mechanism provides a qualitative understanding of the interaction between polymer-chain segments and nanofillers: the nanofillers may disrupt the polymer-chain packing and increase the free volume between the polymer chains, enhancing gas diffusion and, in turn, increas-

ing gas permeability. This mechanism is consistent with many experimental observations [20,22,23,48–50]. 5.3. Solubility increase mechanism The solubility increase mechanism is based on the interaction between the penetrants and the nanofillers. Functional groups, such as hydroxyl, on the surface of the inorganic nanofiller phase may interact with polar gases, such as CO2 and SO2 , and increase the penetrants’ solubility in the nanocomposite membranes and, in turn, increase the gas permeability. For example, in the nanocomposite membranes of PEBAX/silica, Kim and Lee [26] reported that the high CO2 permeability and CO2 /N2 permselectivity increases of the nanocomposite membranes arose from the strong interaction between CO2 molecules and the residual hydroxyl groups on the silica domain, and the additional sorption sites in polyamide block of PEBAX. In the nanocomposite membranes of 6FPAI/TiO2 , Hu et al. [51] also observed a strong interaction between the CO2 and TiO2 domains. The gas permeability (P) for the 6FPAI and 6FPAI/TiO2 membranes were analyzed by using the Arrhenius equation [81,82]:   −Ep P = P0 exp (3) ; Ep = Ed + Hs RT where P0 is a preexponential factor, Ep the apparent activation energy equal to the activation energy of diffusion (Ed ) plus the enthalpy of sorption (Hs ), R the ideal gas constant, and T is absolute temperature. The authors concluded that despite of the influence of Ed , the interaction between residual −OH groups on TiO2 and the polar CO2 molecules decreased Hs for CO2 in the 6FPAI/TiO2 nanocomposite membrane as compared to the unfilled 6FPAI membrane, which would decrease Ep and lead to an increase in gas permeability [51]. 5.4. Nanogap hypothesis In our study of nanocomposite membranes of BPPOdp and silica [83], we found that unmodified silica dispersed rather heterogeneously in the membranes, but greatly improved diffusivities and permeabilities of CO2 and CH4 without changing the CO2 /CH4 selectivity compared with the pure BPPOdp membrane (see Table 2). The permeabilities of the gases increased

Table 2 Gas-separation performance of BPPOdp /surface modified 10 nm silica nanocomposite membranes [83] Nanocomposite membranesa

PCO2 (Barrer)

DCO2 × 108b (cm2 /s)

SCO2 c (cm3 (STP)/cm3 )

PCH4 (Barrer)

DCH4 × 108 (cm2 /s)

SCH4 (cm3 (STP)/cm3 )

αCO2 /CH4

Pure BPPOdp BPPOdp /9 wt.% 10 nm silica BPPOdp /9 wt.% trimethylsilyl modified 10 nm silica BPPOdp /9 wt.% triphenylsilyl modified 10 nm silica

78.0 177.0 104.0

8.73 19.7 10.2

9.79 9.80 11.1

5.00 11.6 7.53

2.20 5.60 3.91

2.40 2.30 2.09

15.6 15.3 13.8

112.7

9.75

12.6

7.88

4.36

1.96

14.3

a b c

Test condition: 10 psig feed pressure and room temperature. Diffusivity of CO2 . Solubility of CO2 .

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nomical, high-performance gas-separation membranes because the membrane is relatively easy to prepare and suitable for dispersing all kinds of inorganic materials in the organic matrix. Modification of fillers and matrices has become an expanding field of research as the introduction of functional groups can improve dispersion of fillers and change chemical affinities of penetrants in nanocomposite membranes. Much research and development are still needed to develop polymer–inorganic nanocomposite membranes for gas separation. Acknowledgement We thank Wyoming’s Enhanced Oil Recovery Institute (EORI) for financial support. Fig. 5. Illustration of nanogap formation in the BPPOdp /silica nanocomposite membranes [83].

with increasing silica concentration. For example, the PCO2 of the BPPOdp /10 nm silica membrane was 177 Barrer at 9 wt.% silica, and reached 436 Barrer at 23 wt.% silica, about 5.6 times of that of the pure BPPOdp membrane, but the selectivity over CH4 remained unchanged (see Tables 1 and 2). We thus proposed that the nanoparticles having better compatibility with BPPOdp such as trimethylsilyl- or triphenylsilyl-modified silica, which dispersed more homogeneously in the polymer matrix, would more efficiently disrupt the polymer-chain packing and increase the free volume for molecular diffusion and thus gas permeability. However, as shown in Table 2, these homogeneous nanocomposite membranes with the modified silica nanoparticles had substantially decreased gas permeability compared to the membranes with the unmodified silica mainly because of the decreased gas diffusivity for both CO2 and CH4 . This finding could not be explained well by the chainunpacking-caused free-volume increase mechanism, suggesting that in the BPPOdp /silica composite membrane, the increased gas permeability did not result from the disrupted polymer-chain packing. Accordingly, we proposed that due to the poor compatibility of the silica surface and the polymer, the polymer chains could not tightly contact the silica nanoparticles, thus forming a narrow gap surrounding the silica particles (see Fig. 5). The gas diffusion path was shortened and thus the apparent gas diffusivity and permeability were increased. This also explained why the addition of the nanoparticles enhanced gas permeability but did not affect the gas selectivity. Once the nanoparticle surface was compatible with the polymer, the nanogaps could not form any more due to the tight contact between the polymer and the filler particles. In another study of BPPOdp /CNT and BPPOdp /COOH-CNT nanocomposite membranes [65], we found that the nanogap hypothesis also explained those experiment results very well. Additionally, Moore and Koros [84] also observed the generation of interface gaps in the Udel® polymer/zeolite 4A composite membranes. 6. Conclusions and future directions Nanocomposite membranes with inorganic nanofillers embedded in a polymer matrix have potentials to provide eco-

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