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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

bstract

Polymer–inorganic nanocomposite membranes present an interesting approach to improve the separation properties of polymer membranesecause they possess properties of both organic and inorganic membranes such as good permeability, selectivity, mechanical strength, and thermalnd chemical stability. The preparations and structures of polymer–inorganic nanocomposite membranes, their applicability to gas separation andeparation mechanism are reviewed. 2007 Elsevier B.V. All rights reserved.

eywords: Nanocomposite membranes; Polymer membrane; Gas separation; Gas transport mechanism

ontents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2812. Types of nanocomposite membrane by structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2823. Preparation of nanocomposite membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283

3.1. Solution blending . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2833.2. In situ polymerization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2833.3. Sol–gel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 283

4. Gas separation properties of nanocomposite membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2835. Gas transport mechanisms in nanocomposite membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 287

5.1. Maxwell’s model . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2875.2. Free-volume increase mechanism . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 287

5.3. Solubility increase mechanism . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2885.4. Nanogap hypothesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 288

6. Conclusions and future directions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. . . .. . . .

Acknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

. Introduction

In the last two decades significant improvements in the per-ormance of polymeric membranes for gas separation have

een made [1–6], and our understanding of the relationshipsmong the structure, permeability and selectivity of polymericembranes has been greatly advanced [2,3]. Newer polymericembrane materials such as polyimides (PI) and cross-linked

∗ Corresponding author. Tel.: +1 307 7662468; fax: +1 307 7666777.E-mail address: sheny@uwyo.edu (Y. Shen).

383-5866/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.seppur.2006.12.017

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 289

polyethylene glycol (PEG) have been continuously developed[7–11]. Some polymeric membranes have already been used inindustry [12,13]. For instance, a plant separating air into its con-stituent gases and producing pure nitrogen at nearly 24 t h−1 inBelgium by Praxair Co. began operation in 1996 [3]. Polymericmembranes tend to be more economical than other membranesbecause of their ability to be easily spun into hollow fibersor spiral-wound modules due to their flexibility and solution

processability [1].

Despite these advantages and progresses, polymeric mem-branes are still restricted by the trade-off trend between gaspermeability and selectivity, as suggested by Robeson [14].

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

Nomenclature

A preexponential factorAF2400 poly(2,2-bis(trifluoromethyl)-4,5-difluoro-1,3-

dioxole-co-tetrafluoroethylene)AIBN 2,2′-azobisisobutyronitrileAPrMDEOS aminopropylmethyldiethoxysilaneAPrTMOS aminopropyltrimethoxysilaneBPPOdm poly(2,6-dimethyl-1,4-phenylene oxide)BPPOdp poly(2,6-diphenyl-1,4-phenylene oxide)C60 fullereneCNTs carbon nanotubesCOOH-CNTs carboxylic acid-functionalized CNTsD diffusivity or diffusion coefficientsDA diffusivity of gas ADABA diaminobenzoic acidDSC differential scanning calorimetryEd energy of diffusionEp apparent activation energy6FDA hexafluoroisopropylidene diphthalic anhydrideFPAI fluorinated poly(amide-imide)6FpDA hexafluoroisopropylidene dianiline�Hs enthalpy of sorptionMA methacrylic acidMTMOS methyltrimethoxysilaneNMP N-methylpyrrolidoneP permeabilityPA permeability of gas APc gas permeability in composite membranePp gas permeability in pure polymer matrixP0 preexponential factorPA polyamidePAA polyamic acidPAI poly(amide-imide)PALS position annihilation lifetime spectroscopyPAN polyacrylonitrilePEBAX poly(amide-6-b-ethylene oxide)PEG polyethylene glycolPEI poly(ether imide)PEO poly(ethylene oxide)PI polyimidesPMA poly(methacrylic acid)PMDA pyromellitic dianhydridePMP poly(4-methyl-2-pentyne)PPEPG PPG-block-PEG-block-PPGPPG poly(propylene glycol)PSF polysulfonePTMOS phenyltrimethoxysilanePTMSP poly(1-trimethylsilyl-1-propyne)R ideal gas constantS solubilitySA solubility of gas ASEM scanning electron microscopyT absolute temperatureTg glass transition temperatureTEM transmission electron microscopy

TEOS tetraethoxysilaneTMOS tetramethoxysilaneVf average free volumeV* minimum free volume element sizeWAXD wide-angle X-ray diffractionXRD X-ray diffraction

Greek symbolsαA/B permeability selectivity of gas A to B

Mltonsmragidmel

aattpmsbacombotmtscma

2

b(

γ overlap factorΦf volume fraction of nanofiller in the membrane

odifications of the chemical structure of a polymer oftenead to an improvement in permeability at the cost of selec-ivity, or vice versa [3]. Additionally, the segmental flexibilityf polymeric membranes often limits their ability to discrimi-ate similar-sized penetrants and they often lose performancetability at high temperatures [1]. On the other hand, inorganicembrane materials such as molecular sieving materials usually

ely on a difference in molecular size to achieve separation. Onlaboratory scale, these membranes show extremely attractiveas permeation and separation performance [15,16]. However,t is still difficult and expensive to fabricate large membranesue to their fragile structures [1,2,17]. Therefore, polymericembranes are still attractive, but alternate approaches that can

nhance their gas separation properties well above the Robesonine are needed.

Polymer–inorganic nanocomposite materials, herein defineds inorganic nanofillers dispersed at a nanometer level in

polymer matrix, have been investigated for gas separa-ion, and have the potential to provide a solution to therade-off problem of polymeric membranes [18,19]. For exam-le, many polymer–inorganic nanocomposite membranes showuch higher gas permeabilities but similar or even improved gas

electivities compared to the corresponding pure polymer mem-ranes [20–26]. The nanocomposite materials may combine thedvantages of each material: for instance, the flexibility and pro-essability of polymers, and the selectivity and thermal stabilityf the inorganic fillers. Additionally, the gas separation perfor-ance of nanocomposite membranes can be further enhanced

y chemical modification [27]. For instance, the introduction ofrganic functional groups on an inorganic filler surface some-imes contributes to not only a better dispersion of the inorganic

aterial in the polymer membrane, but also a better absorp-ion and transportation of penetrants, which results in favorableelectivity and permeability [27,28]. Membrane structure can beontrolled by either the degree of cross-linking of the polymeratrix, or the types of connection bonds between the polymer

nd inorganic phases in the nanocomposite material [28,29].

. Types of nanocomposite membrane by structure

As shown in Fig. 1, polymer–inorganic nanocomposite mem-ranes 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 Purifica

Fig. 1. Illustration of different types of polymer–inorganic nanocomposite mem-bpb

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ranes. (a) Polymer and inorganic phases connected by covalent bonds and (b)olymer and inorganic phases connected by van der Waals force or hydrogenonds.

nd (b) polymer and inorganic phases connected by van deraals force or hydrogen bonds [30].

. Preparation of nanocomposite membranes

Because of the huge difference between the polymer and inor-anic materials in their properties and strong aggregation of theanofillers, polymer–inorganic nanocomposite membranes can-ot be prepared by common methods such as melt blending andoller mixing. The most commonly used preparation technolo-ies for the fabrication of nanocomposite membranes can beivided into the following three types [31].

.1. Solution blending

Solution blending is a simple way to fabricate polymer–norganic nanocomposite membranes. A polymer is first dis-olved in a solvent to form a solution, and then inorganicanoparticles are added into the solution and dispersed bytirring. The nanocomposite membrane is cast by removinghe solvent. For example, Genne et al. [32] prepared polysul-one (PSF)/ZrO2 nanocomposite membranes using 18 wt.% PSFolution in N-methylpyrrolidone (NMP) with adding variousmounts of ZrO2 nanoparticles. The membrane permeabilityncreased as the ZrO2 weight fraction increased. Wara et al.33] reported the fabrication of nanocomposite membranes ofellulose/Al2O3 by using the solution blending.

The solution blending method is easy to operate and suitableor all kinds of inorganic materials, and the concentrations ofhe polymer and inorganic components are easy to control [34].owever, the inorganic ingredients are liable to aggregate in theembranes [35,36].

.2. In situ polymerization

In this method, the nanoparticles are mixed well with organiconomers, and then the monomers are polymerized. There are

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tion Technology 55 (2007) 281–291 283

ften some functional groups such as hydroxyl, carboxyl onhe surface of inorganic particles, which can generate initiatingadicals, cations or anions under high-energy radiation, plasmar other circumstances to initiate the polymerization of theonomers on their surface. For instance, nanocomposite mem-

ranes of poly(methacrylic acid) (PMA)/TiO2 were synthesizedrom TiO2 nanopowder/methacrylic acid dispersions undericrowave radiation [37]. Doucoure et al. [38] reported in situ

lasma polymerization of fluorinated monomers on mesoporousilica membranes. Patel et al. [39,40] prepared cross-linkedanocomposite membranes of PEG/silica and poly(propylenelycol) (PPG)/silica by dispersing silica nanoparticles iniacrylate-terminated PEG and PPG, and subsequent radicalolymerization initiated by 2,2′-azobisisobutyronitrile (AIBN).unes et al. [41] reported the fabrication of nanocompositeembranes of poly(ether imide) (PEI)/SiO2 by using in situ

olymerization.In the in situ polymerization method, inorganic nanoparticles

ith functional groups can be connected with polymer chains byovalent bonds. However, it is still difficult to avoid the aggre-ation of inorganic nanoparticles in the formed membranes.

.3. Sol–gel

The sol–gel method is the most widely used preparationechnology for nanocomposite membranes. In this method,rganic monomers, oligomers or polymers and inorganicanoparticle precursors are mixed together in the solution.he inorganic precursors then hydrolyze and condense intoell-dispersed nanoparticles in the polymer matrix. The advan-

age of this method is obvious: the reaction conditions areoderate—usually room temperature and ambient pressure, and

he concentrations of organic and inorganic components are easyo control in the solution. Additionally, the organic and inorganicngredients are dispersed at the molecular or nanometer leveln the membranes, and thus the membranes are homogeneous42–45].

For example, Iwata et al. [46] reported that by using theol–gel method, a nanocomposite membrane of polyacryloni-rile (PAN) with hydrolysate of tetraethoxysilane (TEOS) as thenorganic phase showed a good performance in O2/N2 sepa-ation. Gomes et al. [47] prepared nanocomposite membranesf poly(1-trimethylsilyl-1-propyne) (PTMSP)/silica by sol–gelopolymerization of TEOS with different organoalkoxysilanesn the tetrahydrofuran solution of PTMSP. Fig. 2 shows the scan-ing electron microscopy (SEM) images of cross-sections ofTMSP membrane (a) and PTMSP/silica membrane (b) pre-ared by sol–gel process.

. Gas separation properties of nanocompositeembranes

The permeability (P) of a gas through a membrane is pro-

ortional to the solubility (S) and diffusivity (D) of the gas inhe membrane (P = D × S). Thus, adding inorganic nanofillers

ay affect the gas separation in two ways: the interactionetween polymer-chain segments and nanofillers may disrupt

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

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ig. 2. SEM photomicrographs of cross-sections of: (a) PTMSP membrane aReproduced with permission from Elsevier Co.)

he polymer-chain packing and increase the voids (free volumes)etween the polymer chains, and thus enhance gas diffusion20,22,23,48–50]; the hydroxyl and other functional groups onhe surface of the inorganic phase may interact with polar gasesuch as CO2 and SO2, improving the penetrants’ solubility inhe nanocomposite membranes [26,51]. Various combinations ofolymers and nanofillers have been tested for gas separations.he results are summarized in Table 1.

Among these nanocomposite membranes, polyimide/silicaaterials have received the most attention for the gas permeation

tudies [28,52,53]. Joly et al. [21,54] fabricated the poly-mide/silica membranes containing 32 wt.% silica via the sol–gel

ethod by adding tetramethoxysilane (TMOS) to polyamic acidPAA) solution and subsequently imidizing at 60–300 ◦C. Ashown in Table 1, the nanocomposite membrane had a higherermeability for CO2 (PCO2 = 2.8 Barrer) and a greater CO2/N2electivity (αCO2/N2 = 22) compared to the polyimide mem-rane (PCO2 = 1.8 Barrer; αCO2/N2 = 18). The gas permeationesults were analyzed using the dual sorption model. In thisodel, it was assumed that the gas molecules dissolved in the

olymer could be classified into two distinct populations: (a)enry-type dissolution and (b) Langmuir-type sorption. The

uthors attributed the increased gas permeation of the nanocom-osite membrane compared to the polyimide membrane tonhanced gas solubility due to an increased contribution ofenry’s type dissolution. Using X-ray diffraction (XRD) andEM, the authors showed that the addition of TMOS to theAA induced some morphological modifications in the polymeratrix.Kusakabe et al. [55] reported that the CO2 permeability in a

olyimide/SiO2 hybrid nanocomposite membrane was 15 timesager than that in the corresponding polyimide. The permselec-ivity of CO2 to N2 was 25 at 30 ◦C. Contributions of the silicand polyimide phases to the composite membrane’s permeance

Tdda

) PTMSP/silica nanocomposite membrane prepared by sol–gel process [47].

ere analyzed using a two-phase permeation model. The effec-ive thickness of the rate-controlling polyimide phase was lesshan one tenth of the thickness of the composite membrane.

Homogeneous nanocomposite membranes of polyimide–iloxane copolymers containing different silica contents wererepared by Smaihi et al. [56] via the sol–gel process of pyromel-itic dianhydride (PMDA), aminoalkoxysilane, and TMOS.hey used two coupling agents, aminopropyltrimethoxysi-

ane (APrTMOS) and aminopropylmethyldiethoxysilaneAPrMDEOS) to provide bonding between the imide andhe inorganic silica. Higher gas permeability was observedor the membrane using APrMDEOS than for the membranesing APrTMOS at the same silica content. IR studies of theanocomposite material revealed that the presence of methylide groups linked to the silicon of the APrMDEOS precursornhibited the formation of −OH linked bonds in the material.

Moaddeb and Koros [22] studied the gas transportationroperties of thin polyimide membranes in the presence ofilica particles. In the nanocomposite membranes on silica-mpregnated aluminum oxide substrates, the presence of silicamproved the gas separation properties of the polyimide layer,articularly for O2 and N2. The increase in permeability was dueo silica-disrupting the polymer-chain packing. The observedignificant increase in the glass-transition temperature (Tg) sug-ested the restriction of chain segmental mobility possibly dueo adsorption of the polymer to the silica surface.

Cornelius and co-workers [23,57–59] studied the effectsf alkoxysilanes, their loading, and the morphology of theesulting polyimide/silica nanocomposite membranes on theermselectivity of several gases including CO2, N2, and CH4.

he polyimides were prepared from hexafluoroisopropyli-ene diphthalic anhydride (6FDA)–hexafluoroisopropylideneianiline (6FpDA)–diaminobenzoic acid (DABA) and thelkoxysilanes included phenyltrimethoxysilane (PTMOS),

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tion Technology 55 (2007) 281–291 285

ethyltrimethoxysilane (MTMOS), TMOS, and TEOS. Theseanocomposite membranes were annealed at 400 ◦C to drivehe sol–gel reactions to a greater extent. In general, the anneal-ng process increased the gas permeation of the nanocomposite

embranes by about 200–500%, while the permselectivityropped slightly. An exception was the 6FDA–6FpDA–DABA-5 (containing 25 mol% DABA) membrane. The nanocompositeembranes with 22.5 wt.% TMOS and MTMOS both had

ncreased CO2 permeabilities (see Table 1) and CO2/CH4ermselectivity [23]. The authors attributed the increase inas permeation to changes in the free volume distribution andnhanced local segment mobility of the chain ends resulted fromhe removal of sol–gel condensation and polymer degradationy-products.

Suzuki and Yamada [24] reported the physical and gas trans-ort properties of a 6FDA-based hyperbranched polyimide/silicaanocomposite membrane prepared using polyamic acid, waternd TMOS via the sol–gel technique. CO2, O2 and N2 per-eability coefficients of the membrane increased as the silica

ontent increased. It was pointed out that the increased gas per-eabilities were mainly attributable to the increase in the gas

olubilities. In contrast, CH4 permeability of the nanocompositeembranes decreased with increasing silica content because of

he decrease in the CH4 diffusivity. As a result, CO2/CH4 selec-ivity (see Table 1) of the nanocomposite membranes increasedemarkably. This kind of nanocomposite membrane had highhermal stability and excellent gas selectivity, and is expected toe a high-performance gas-separation membrane.

Polymer–inorganic nanocomposite membranes based on auorinated poly(amide-imide) (FPAI) and TiO2 were fabri-ated by Hu et al. [51] via the sol–gel method. An aromaticoly(amide-imide) (PAI) was chosen as the polymer matrixaterial because it provided superior mechanical properties,

igh thermal stability, solvent resistance, and high permeability.he nanocomposite membrane had a more rigid and dense struc-

ure than the corresponding pure FPAI membrane. The authorsbserved a specific interaction between such gases as CO2 and2 and the TiO2 particles in the nanocomposite membrane.igher selectivities for CO2/CH4 and H2/CH4 gas pairs (seeable 1) were observed in the composite membrane containinglow concentration (7.3 wt.%) of TiO2.

Kong et al. [25] prepared polyimide/TiO2 nanocompositeembranes by blending TiO2 sol and a polyimide solution.ecause of the improved TiO2-sol preparation process andlending method, the TiO2 content could reach about 40 wt.%n the membrane. There existed a strong interaction betweenhe TiO2 phase and the polyimide phase. A higher TiO2 contentn the membrane resulted in a greater enhancement of the gas-eparation performance. The H2 and O2 permeabilities of theembrane with 25 wt.% TiO2 were 14.1 and 0.72 Barrer, respec-

ively, which were 3.7 times and 4.3 times higher than thosef the pure polyimide. The selectivities of H2/N2 and O2/N2see Table 1) were also slightly improved compared to the pure

olymer membrane (αH2/N2 = 167, αO2/N2 = 9.3).

Polymer–inorganic nanocomposite membranes of polyamide-6-b-ethylene oxide) (PEBAX)/silica were prepared byim and Lee [26] via in situ polymerization of TEOS using

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86 H. Cong et al. / Separation and Pur

he sol–gel process. The membrane containing 27 wt.% silicaad a CO2 permeability of 277 Barrer and CO2/N2 selectivityf 79, both higher than those of the pure polymer membrane.he PEBAX copolymer consisted of two distinct regions, the

mpermeable crystalline polyamide (PA) phase, and the perme-ble amorphous poly(ethylene oxide) (PEO) phase [60]. Usingide-angle X-ray diffraction (WAXD) and differential scan-ing calorimetry (DSC), the authors showed that the presencef silica domains in the nanocomposite membrane significantlyecreased the degree of crystallinity of the PA phase and causedeorientation of the PEO phase. They reported that the nanocom-osite membranes exhibited higher gas permeability coefficientsnd permselectivities than the PEBAX alone, particularly at anlevated temperature. The gas permeability and permselectivityncreased as the silica content increased in the nanocomposite

embrane. It was concluded that the increases in permeabilitynd permselectivity of the nanocomposite membrane arose fromhe strong interaction between CO2 molecules and the residualydroxyl groups on the silica domain, and the additional sorptionites in polyamide block of PEBAX.

Higuchi et al. [48] observed an increase in gas perme-bility upon adding fullerene (C60) particles to polystyrene.ermeability increased 41% for ethylene, 47% for N2 and 75%or ethane in a film containing 10 wt.% fullerene particles at5 ◦C. Selectivity decreased about 25% for O2/N2 (see Table 1)nd ethylene/ethane. The enhancement in permeability wasttributed to increases in diffusion coefficients caused by thencreased free volume in the membrane.

Pinnau and He [61,62] and Merkel et al. [20,63] reportedhat adding nanosized impermeable particles of commercialumed silica to poly(4-methyl-2-pentyne) (PMP) increased gasnd vapor permeabilities with increased particle loading. Forxample, the n-butane permeability (see Table 1) increased by aactor of 3 relative to that of the pure PMP when 30 wt.% silicaas added at 25 ◦C. Silica particles did not alter the solubil-

ty of the nanocomposite, but they did significantly increase theas diffusion coefficient. For instance, the CH4 diffusion coeffi-ient doubled at 30 wt.% silica loading [63]. The silica particles

∼12 nm) used in these studies were small enough to disruptolymer-chain packing in the polymers, which resulted in anncrease in polymer fractional free volume [20]. The free volumencrease was characterized using density and position annihila-

Citg

Fig. 3. TEM photomicrographs of cross-sections of BPPOdp/single-wall CN

on Technology 55 (2007) 281–291

ion lifetime spectroscopy (PALS) measurements. By increasingractional free volume, gas diffusion coefficients and, in turn,as permeability increased [49,64]. Particle loading influencedhe gas transport properties in these nanocomposite membranes.or example, N2 permeability tripled as silica loadings in PMP

ncreased from 5 to 25 vol.% [62]. The enhancement in n-butaneermeability coincided with a substantial enhancement in n-utane/CH4 mixed gas selectivity. The n-butane/CH4 selectivitysee Table 1) doubled relative to that of the pure PMP at a 30 wt.%ilica concentration. The enhanced n-butane permeability and-butane/CH4 selectivity were attributed to the increased freeolume of PMP by the silica-particles [62].

Recently, our group fabricated nanocomposite membranes ofrominated poly(2,6-diphenyl-1,4-phenylene oxide) (BPPOdp)nd carbon nanotubes (CNTs) via the solution-blending method65]. CNTs were chosen as the inorganic filler material becausehey are very effective in reinforcing polymeric materialsnd thus may provide better mechanical properties [66–68]o the membrane. Fig. 3 shows the transmission electron

icroscopy (TEM) photomicrographs of cross-sections ofPPOdp/single-wall CNT nanocomposite membranes preparedy solution-blending method. The composite membranes hadn increased CO2 permeability but a similar CO2/N2 selectiv-ty (see Table 1) compared to the corresponding pure BPPOdpembranes (PCO2 = 78 Barrer; αCO2/N2 = 30). The CO2 per-eability increased with increasing the carbon nanotube content

nd reached a maximum of 155 Barrer at 9 wt.% single-wallNTs, or 148 Barrer at 5 wt.% multiwall CNTs. The CO2/N2

eparation performance was not sensitive to the CNT diam-ter and length. However, the carboxylic acid-functionalizedNTs (COOH-CNTs), which were more homogeneously dis-ersed in BPPOdp, neither increased the gas permeability noreteriorated the gas separation performance. We concluded thatue to the incompatibility of pristine CNTs and the BPPOdphains, the polymer chains did not attach to the CNT wall tightly,orming narrow gaps surrounding the CNTs. Gas moleculeshus easily passed through the gap and had a shortcut. Thislso explained why the addition of CNTs did not affect the

O2/N2 selectivity. The modified CNTs surface was compat-

ble with the polymer and the polymer chains could packightly on the CNT surface, closing the nanogap and leaving theas permeability unenhanced. In another study, nanocompos-

T nanocomposite membranes prepared by solution-blending method.

rifica

iosiCpiP

9fim

ptinfmhsnwrcmhcsosms(b(iPpo

5m

5

mMtus

P

wtf

ddiaf

lfimAaMtnnn(Pp

itbt

5

sdCexpression for penetrant diffusion coefficients (D) [63]:

D = A exp

(−γV ∗

Vf

)(2)

H. Cong et al. / Separation and Pu

te membranes of brominated poly(2,6-dimethyl-1,4-phenylenexide) (BPPOdm) and silicas were fabricated successfully via theolution-blending method [69], and we also found the compos-te membranes had an increased CO2 permeability but a similarO2/N2 selectivity (see Table 1) compared to the correspondingure BPPOdm membranes. The permeabilities of all the gasesncreased with increasing silica concentration. For example, theCO2 of the BPPOdm/10 nm silica membrane was 187 Barrer atwt.% silica, and reached 523 Barrer at 23 wt.% silica, aboutve times of that of the pure BPPOdm membrane; in the sameembranes the selectivity over N2 remained unchanged.The gas-separation performance of the present nanocom-

osite membranes can be further enhanced by modification ofhe fillers and matrices. For example, Patel et al. [27] stud-ed the effects of nanoparticle functionality on CO2-selectiveanocomposite membranes derived from cross-linked PEG andound that methacrylate-functionalized silica nanoparticles wereore effective in improving rheological properties and retaining

igh CO2 selectivity than the original hydroxyl-functionalizedilica nanoparticles of comparable size in the cross-linkedanocomposite membranes. The reason for this differenceas that the methacrylate-functionalized silica nanoparticles

eacted with PEG-diacrylate oligomers in the cross-link pro-ess, and thus improved dispersion of fillers in the polymeratrix and increased the interaction between them, while the

ydroxyl-functionalized silica nanoparticles could not attend theross-link reaction. Kim et al. [70] reported that by using theol–gel method, the nanocomposite membrane with PEG as therganic phase and hydrolysate of TEOS as the inorganic phasehowed good performance in CO2/N2 separation. The CO2 per-eability was 94.2 Barrer with a CO2/N2 selectivity of 38.3. By

imply changing the PEG matrix to PPG-block-PEG-block-PPGPPEPG), Sforca et al. [71] prepared nanocomposite membranesy the same method and reported that the CO2 permeabilitysee Table 1) increased to 125 Barrer, and the CO2/N2 selectiv-ty increased to 89. It was concluded that the introduction ofPG segments in the PEG chains not only disrupted the originalolymer-chain packing but also changed the chemical affinitiesf penetrants in the matrix.

. Gas transport mechanisms in nanocompositeembranes

.1. Maxwell’s model

Adding impermeable inorganic nanoparticles to a poly-er is typically expected to reduce the gas permeability [72].axwell’s model, developed to analyze the steady-state dielec-

ric properties of a diluted suspension of spheres [73], is oftensed to model permeability in membranes filled with roughlypherical impermeable particles [74]:

c = Pp

(1 − Φf

)(1)

1 + 0.5Φf

here Pc and Pp are the permeability of the nanocomposite andhe pure polymer matrix, respectively, and Φf is the volumeraction of the nanofiller.

F2C

tion Technology 55 (2007) 281–291 287

The numerator represents the loss of membrane solubilityue to the loss of polymer volume available for sorption. Theenominator represents a decrease in diffusivity due to increas-ng the penetrant diffusion pathway length [63]. Both factorsct to decrease permeability with increasing particle volumeraction [75].

Maxwell’s model partly explains the gas permeabilityoss in some nanocomposite membranes, especially in theullerene-particle-filled polymer membranes [76–78]. However,n general, the addition of fullerene to polymers decreases per-

eability more than the loss predicted by Maxwell’s model [75].lso, Higuchi et al. [48] observed an increase in gas perme-

bility upon adding fullerene particles to polystyrene, whereasaxwell’s model predicted a 14% loss in gas permeability for

his system. Additionally, more and more polymer–inorganicanocomposite membranes have been observed with similaron-Maxwellian effects [20,22,23,62,65]. For example, addinganosized impermeable particles of commercial fumed silicaTS 530, 12 nm primary diameter) to a glassy polymer (e.g.,MP) increased gas and vapor permeabilities with increasedarticle loading (see Fig. 4) [62].

The problem with Maxwell’s model lies in its neglect of thenteractions between the nanofillers and the polymer chains, andhe nanofillers and the penetrants. In most nanocomposite mem-ranes, these interactions are strong, and significantly changehe diffusivity and solubility of penetrants.

.2. Free-volume increase mechanism

The effect of polymer free volume on penetrant diffu-ion coefficients is often modeled by the statistical-mechanicalescription of diffusion in a liquid of hard spheres proposed byohen and Turnbull [79]. This model provides the following

ig. 4. N2 permeability enhancement of PMP as a function of filler content at5 ◦C and 50 psig feed pressure [62]. (Reproduced with permission from Elseviero.)

2 ificati

wafsaite

rs4emsmmovsiaaiainisrvimCcasp

tspp

ie

5

bsmti

Ppnbopobfu

P

weeaiotut

5

s

TG

N

PBB

B

88 H. Cong et al. / Separation and Pur

here A is a preexponential factor weakly dependent on temper-ture, γ an overlap factor introduced to avoid double-countingree volume elements, V* the minimum free volume elementize that can accommodate a penetrant molecule (and is closelyssociated with penetrant size), and Vf is the average free volumen the media accessible to penetrants for transport. Accordingo Eq. (2), an increase in polymer free volumes is expected tonhance penetrant diffusion.

Based on the PALS measurements, Merkel et al. [49]eported that the addition of fumed silica increased theize of free-volume elements in poly(2,2-bis(trifluoromethyl)-,5-difluoro-1,3-dioxole-co-tetrafluoroethylene) (AF2400). Thenhanced free volume of AF2400/silica resulted in aug-ented penetrant permeability and diffusion coefficients,

imilar to the observation in the PMP/silica nanocompositeembranes [20,63]. The authors concluded that the improve-ent in permeability reflected hybridization-induced disruption

f polymer-chain packing and an accompanying elevated freeolume available for molecular diffusion. Winberg et al. [50]tudied the free volume in silica-filled PTMSP nanocompos-te membranes with PALS at filler concentrations between 0nd 50 wt.%. A bimodal free-volume distribution was observed,nd the size of larger free volume cavities was significantlyncreased upon addition of hydrophobic fumed silica. Theuthors observed a strong correlation between N2 permeabil-ty and the volume of the larger free-volume cavities in theanocomposite membranes, and the permeability increased withncreasing filler content. It is worth to mention that Hill [80]ignificantly extended the Cohen–Turnbull free volume theoryecently by hypothesis that the accompanying increase in freeolume in the nanocomposite membranes reflects a repulsiventeraction between the polymer chains and inclusions during

embrane casting. He proposed a theoretical model based on theohen–Turnbull statistical mechanical theory, which not onlyaptured the correct dependence of the diffusive permeabilitynd selectivity of polymeric nanocomposites on the inclusionize and volume fraction, but also achieved a quantitative inter-retation of the Merkel’s experiments [20,63].

The free-volume increase mechanism provides a qualita-

ive understanding of the interaction between polymer-chainegments and nanofillers: the nanofillers may disrupt theolymer-chain packing and increase the free volume between theolymer chains, enhancing gas diffusion and, in turn, increas-

hftb

able 2as-separation performance of BPPOdp/surface modified 10 nm silica nanocomposit

anocomposite membranesa PCO2 (Barrer) DCO2 × 108b

(cm2/s)SCO2

c

(cm3 (STP)

ure BPPOdp 78.0 8.73 9.79PPOdp/9 wt.% 10 nm silica 177.0 19.7 9.80PPOdp/9 wt.% trimethylsilylmodified 10 nm silica

104.0 10.2 11.1

PPOdp/9 wt.% triphenylsilylmodified 10 nm silica

112.7 9.75 12.6

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

on Technology 55 (2007) 281–291

ng gas permeability. This mechanism is consistent with manyxperimental observations [20,22,23,48–50].

.3. Solubility increase mechanism

The solubility increase mechanism is based on the interactionetween the penetrants and the nanofillers. Functional groups,uch as hydroxyl, on the surface of the inorganic nanofiller phaseay interact with polar gases, such as CO2 and SO2, and increase

he penetrants’ solubility in the nanocomposite membranes and,n turn, increase the gas permeability.

For example, in the nanocomposite membranes ofEBAX/silica, Kim and Lee [26] reported that the high CO2ermeability and CO2/N2 permselectivity increases of theanocomposite membranes arose from the strong interactionetween CO2 molecules and the residual hydroxyl groupsn the silica domain, and the additional sorption sites inolyamide block of PEBAX. In the nanocomposite membranesf 6FPAI/TiO2, Hu et al. [51] also observed a strong interactionetween the CO2 and TiO2 domains. The gas permeability (P)or the 6FPAI and 6FPAI/TiO2 membranes were analyzed bysing the Arrhenius equation [81,82]:

= P0 exp

(−Ep

RT

); Ep = Ed + �Hs (3)

here P0 is a preexponential factor, Ep the apparent activationnergy equal to the activation energy of diffusion (Ed) plus thenthalpy of sorption (�Hs), R the ideal gas constant, and T isbsolute temperature. The authors concluded that despite of thenfluence of Ed, the interaction between residual −OH groupsn TiO2 and the polar CO2 molecules decreased �Hs for CO2 inhe 6FPAI/TiO2 nanocomposite membrane as compared to thenfilled 6FPAI membrane, which would decrease Ep and leado an increase in gas permeability [51].

.4. Nanogap hypothesis

In our study of nanocomposite membranes of BPPOdp andilica [83], we found that unmodified silica dispersed rather

eterogeneously in the membranes, but greatly improved dif-usivities and permeabilities of CO2 and CH4 without changinghe CO2/CH4 selectivity compared with the pure BPPOdp mem-rane (see Table 2). The permeabilities of the gases increased

e membranes [83]

/cm3)PCH4

(Barrer)DCH4 × 108

(cm2/s)SCH4

(cm3 (STP)/cm3)αCO2/CH4

5.00 2.20 2.40 15.611.6 5.60 2.30 15.37.53 3.91 2.09 13.8

7.88 4.36 1.96 14.3

H. Cong et al. / Separation and Purifica

Fm

wtsoCpBwwipnctd

utgpicagfwisnmafiap

6

e

ntpMictan

A

(

R

[

[

[

[

[

[

[

[

[

[

ig. 5. Illustration of nanogap formation in the BPPOdp/silica nanocompositeembranes [83].

ith increasing silica concentration. For example, the PCO2 ofhe BPPOdp/10 nm silica membrane was 177 Barrer at 9 wt.%ilica, and reached 436 Barrer at 23 wt.% silica, about 5.6 timesf that of the pure BPPOdp membrane, but the selectivity overH4 remained unchanged (see Tables 1 and 2). We thus pro-osed that the nanoparticles having better compatibility withPPOdp such as trimethylsilyl- or triphenylsilyl-modified silica,hich dispersed more homogeneously in the polymer matrix,ould more efficiently disrupt the polymer-chain packing and

ncrease the free volume for molecular diffusion and thus gasermeability. However, as shown in Table 2, these homogeneousanocomposite membranes with the modified silica nanoparti-les had substantially decreased gas permeability compared tohe membranes with the unmodified silica mainly because of theecreased gas diffusivity for both CO2 and CH4.

This finding could not be explained well by the chain-npacking-caused free-volume increase mechanism, suggestinghat in the BPPOdp/silica composite membrane, the increasedas permeability did not result from the disrupted polymer-chainacking. Accordingly, we proposed that due to the poor compat-bility of the silica surface and the polymer, the polymer chainsould not tightly contact the silica nanoparticles, thus formingnarrow gap surrounding the silica particles (see Fig. 5). The

as diffusion path was shortened and thus the apparent gas dif-usivity and permeability were increased. This also explainedhy the addition of the nanoparticles enhanced gas permeabil-

ty but did not affect the gas selectivity. Once the nanoparticleurface was compatible with the polymer, the nanogaps couldot form any more due to the tight contact between the poly-er and the filler particles. In another study of BPPOdp/CNT

nd BPPOdp/COOH-CNT nanocomposite membranes [65], weound that the nanogap hypothesis also explained those exper-ment results very well. Additionally, Moore and Koros [84]lso observed the generation of interface gaps in the Udel®

olymer/zeolite 4A composite membranes.

. Conclusions and future directions

Nanocomposite membranes with inorganic nanofillersmbedded in a polymer matrix have potentials to provide eco-

[

tion Technology 55 (2007) 281–291 289

omical, high-performance gas-separation membranes becausehe membrane is relatively easy to prepare and suitable for dis-ersing all kinds of inorganic materials in the organic matrix.odification of fillers and matrices has become an expand-

ng field of research as the introduction of functional groupsan improve dispersion of fillers and change chemical affini-ies of penetrants in nanocomposite membranes. Much researchnd development are still needed to develop polymer–inorganicanocomposite membranes for gas separation.

cknowledgement

We thank Wyoming’s Enhanced Oil Recovery InstituteEORI) for financial support.

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