thermally rearranged (tr) polymer membranes for co2 separation

14
Journal of Membrane Science 359 (2010) 11–24 Contents lists available at ScienceDirect Journal of Membrane Science journal homepage: www.elsevier.com/locate/memsci Thermally rearranged (TR) polymer membranes for CO 2 separation Ho Bum Park a,, Sang Hoon Han a , Chul Ho Jung a , Young Moo Lee a,∗∗ , Anita J. Hill b a School of Chemical Engineering and WCU Department of Energy Engineering, Hanyang University, Seoul 133-791, Republic of Korea b The Commonwealth Scientific and Industrial Research Organization (CSIRO) Materials Science and Engineering, Private Bag 33, S Clayton, VIC 3169, Australia article info Article history: Received 13 June 2009 Received in revised form 14 September 2009 Accepted 16 September 2009 Available online 14 October 2009 Keywords: Thermal rearrangement High-temperature polymers Microporous polymers Membranes Gas separation abstract The evolution of micropores in polymer membranes helps accelerate mass transport phenomena on a sub-nanoscale, providing significant technological applications for adsorption, separation and storage. Here we report the synthesis and characterization of thermally rearranged (TR) polymer membranes showing unexpected microporous characters that are often observed in microporous inorganic materi- als, by using thermal rearrangement of various aromatic polyimides with semi-rigid chain segments in a solid state. Differing from other superglassy polymers (e.g., poly(1-methylsilyl-1-propyne) (PTMSP)) possessing larges cavities, these TR polymer membranes show fast molecular transport as well as a molecular sieving effect for small gas molecules. Micropore structures and their size distributions can be easily tuned by varying the monomer structures of the precursor polymers (i.e., polyimides with ortho-positioned functional groups, PIOFG) and by using different thermal treatment protocols (e.g., final temperature and thermal dwell time). These TR polymer membranes exhibit excellent gas separation performance, especially in carbon dioxide separations (e.g., CO 2 /CH 4 ), without any paramount plasti- cization effect. The current approach will be useful in an assessment of the achievements of membrane materials science, providing much insight into new class of microporous polymers. © 2009 Published by Elsevier B.V. 1. Introduction Recently, polymer membranes have served as a key element in many useful scientific and technological fields such as water purification [1], gas and vapor separation [2,3], sensors [4] and fuel cells [5]. In these areas, the heart of their functions is fast, selective transport of molecules (e.g., gases, liquids and ions) to save energy and cost. With the advent of a new era, demanding energy-saving and environment-friendly innovative processes, however, conven- tional polymer membranes for gas separation are confronted with some difficulties—poor chemical and thermal stability with rel- atively low mass transport rates and low separation efficiency. However, organic polymer membranes for gas separation still have many advantages over inorganic and metal substances from the viewpoint of manufacture (e.g., cost, handling, and processability for mass production) [6]. To compete effectively with other mate- rials, it is necessary for polymer membranes to achieve both high permeability and selectivity as well as to impart strong environ- mental adaptability. Corresponding author. ∗∗ Corresponding author. Tel.: +82 2 2220 2338; fax: +82 2 2282 0922. E-mail addresses: [email protected] (H.B. Park), [email protected] (Y.M. Lee). Among many membrane-based gas separations, especially CO 2 - selective separation polymeric membranes (e.g., CO 2 /CH 4 , CO 2 /H 2 , and CO 2 /N 2 ) are extensively studied for the application of the refining of natural gas, fermentation gas, water–gas-shift reac- tion in pre-combustion and coal-based power plant flue gas (for post-combustion) [7,8]. Particularly, demonstrations of sweeten- ing of natural gas using cellulose acetate membranes were made as early as the 1960s, and the first industrial plants for CO 2 removal using cellulose acetate membranes were installed in the 1980s. Since then, there has been intense research on polymeric mem- brane materials for CO 2 /CH 4 separation [9,10]. Glassy polymers, in particular aromatic polyimides, whose selectivity is mainly a result of the large differences in the mobility of the penetrants due to the stiffness of the polymer chains, showed high selec- tivity towards CO 2 based on laboratory results with pure gases. Unfortunately, the high values of ideal selectivity from pure-gas experiments were catastrophically reduced when the membranes were tested with CO 2 /CH 4 mixtures and real natural gas stream [11,12]. This difference between pure and mixed-gas selectivity values is due to membrane plasticization. During high pressure nat- ural gas operations, a considerable amount of carbon dioxide and heavy hydrocarbons is sorbed by the polymer, which swells and dilates, with a consequent increase in the mobility of the polymer chains. The higher chain mobility compromises the size-sieving ability of the polymer, reducing its selectivity. As a result of their plasticization-reduced selectivity, current commercial membranes 0376-7388/$ – see front matter © 2009 Published by Elsevier B.V. doi:10.1016/j.memsci.2009.09.037

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Page 1: Thermally rearranged (TR) polymer membranes for CO2 separation

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Journal of Membrane Science 359 (2010) 11–24

Contents lists available at ScienceDirect

Journal of Membrane Science

journa l homepage: www.e lsev ier .com/ locate /memsci

hermally rearranged (TR) polymer membranes for CO2 separation

o Bum Parka,∗, Sang Hoon Hana, Chul Ho Junga, Young Moo Leea,∗∗, Anita J. Hill b

School of Chemical Engineering and WCU Department of Energy Engineering, Hanyang University, Seoul 133-791, Republic of KoreaThe Commonwealth Scientific and Industrial Research Organization (CSIRO) Materials Science and Engineering, Private Bag 33, S Clayton, VIC 3169, Australia

r t i c l e i n f o

rticle history:eceived 13 June 2009eceived in revised form4 September 2009ccepted 16 September 2009vailable online 14 October 2009

eywords:

a b s t r a c t

The evolution of micropores in polymer membranes helps accelerate mass transport phenomena on asub-nanoscale, providing significant technological applications for adsorption, separation and storage.Here we report the synthesis and characterization of thermally rearranged (TR) polymer membranesshowing unexpected microporous characters that are often observed in microporous inorganic materi-als, by using thermal rearrangement of various aromatic polyimides with semi-rigid chain segments ina solid state. Differing from other superglassy polymers (e.g., poly(1-methylsilyl-1-propyne) (PTMSP))possessing larges cavities, these TR polymer membranes show fast molecular transport as well as a

hermal rearrangementigh-temperature polymersicroporous polymersembranesas separation

molecular sieving effect for small gas molecules. Micropore structures and their size distributions canbe easily tuned by varying the monomer structures of the precursor polymers (i.e., polyimides withortho-positioned functional groups, PIOFG) and by using different thermal treatment protocols (e.g., finaltemperature and thermal dwell time). These TR polymer membranes exhibit excellent gas separationperformance, especially in carbon dioxide separations (e.g., CO2/CH4), without any paramount plasti-cization effect. The current approach will be useful in an assessment of the achievements of membrane

ing m

materials science, provid

. Introduction

Recently, polymer membranes have served as a key elementn many useful scientific and technological fields such as waterurification [1], gas and vapor separation [2,3], sensors [4] and fuelells [5]. In these areas, the heart of their functions is fast, selectiveransport of molecules (e.g., gases, liquids and ions) to save energynd cost. With the advent of a new era, demanding energy-savingnd environment-friendly innovative processes, however, conven-ional polymer membranes for gas separation are confronted withome difficulties—poor chemical and thermal stability with rel-tively low mass transport rates and low separation efficiency.owever, organic polymer membranes for gas separation still haveany advantages over inorganic and metal substances from the

iewpoint of manufacture (e.g., cost, handling, and processabilityor mass production) [6]. To compete effectively with other mate-

ials, it is necessary for polymer membranes to achieve both highermeability and selectivity as well as to impart strong environ-ental adaptability.

∗ Corresponding author.∗∗ Corresponding author. Tel.: +82 2 2220 2338; fax: +82 2 2282 0922.

E-mail addresses: [email protected] (H.B. Park), [email protected]. Lee).

376-7388/$ – see front matter © 2009 Published by Elsevier B.V.oi:10.1016/j.memsci.2009.09.037

uch insight into new class of microporous polymers.© 2009 Published by Elsevier B.V.

Among many membrane-based gas separations, especially CO2-selective separation polymeric membranes (e.g., CO2/CH4, CO2/H2,and CO2/N2) are extensively studied for the application of therefining of natural gas, fermentation gas, water–gas-shift reac-tion in pre-combustion and coal-based power plant flue gas (forpost-combustion) [7,8]. Particularly, demonstrations of sweeten-ing of natural gas using cellulose acetate membranes were made asearly as the 1960s, and the first industrial plants for CO2 removalusing cellulose acetate membranes were installed in the 1980s.Since then, there has been intense research on polymeric mem-brane materials for CO2/CH4 separation [9,10]. Glassy polymers,in particular aromatic polyimides, whose selectivity is mainly aresult of the large differences in the mobility of the penetrantsdue to the stiffness of the polymer chains, showed high selec-tivity towards CO2 based on laboratory results with pure gases.Unfortunately, the high values of ideal selectivity from pure-gasexperiments were catastrophically reduced when the membraneswere tested with CO2/CH4 mixtures and real natural gas stream[11,12]. This difference between pure and mixed-gas selectivityvalues is due to membrane plasticization. During high pressure nat-ural gas operations, a considerable amount of carbon dioxide and

heavy hydrocarbons is sorbed by the polymer, which swells anddilates, with a consequent increase in the mobility of the polymerchains. The higher chain mobility compromises the size-sievingability of the polymer, reducing its selectivity. As a result of theirplasticization-reduced selectivity, current commercial membranes
Page 2: Thermally rearranged (TR) polymer membranes for CO2 separation

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or carbon dioxide removal, based on cellulose acetate and poly-mides, can only compete with amine-based systems for small gasow rates (<40 million scfd), which accounts for less than 5% of theotal market [6,11]. Therefore, improved membrane materials forO2 removal from light gases are required.

In addition, more excellent properties such as thermal andhemical stability have been demanded recently in polymer mem-ranes with the ability to be used in various applications andxtreme conditions. Unfortunately, most glassy polymer mem-ranes are difficult to use at high temperatures above 100 ◦Cecause glassy polymers are usually in a non-equilibrium state atoom temperature and at elevated temperatures around or abovelass transition temperature (Tg) the equilibrium process may behortened, resulting in increasing chain mobility and as a resultecreasing gas selectivity. In some applications, high-temperatureolymer membranes are required to operate at above 300 ◦C,owever, there are only a few membrane materials (e.g., polyben-imidazoles and polyimides) suitable for that purpose [13,14].

Aromatic polymers interconnected with heterocyclic rings, suchs polybenzoxazoles (PBO) and polybenzothiazoles (PBT), have aigid-rod structure with high-torsional energy barriers to rota-ion between two individual phenylene-heterocyclic rings [15].hese features are interesting from the point of view of gas sep-ration membranes, since they could lead to large differences inhe mobility of penetrants according to their sizes, and, hence,o high selectivities. However, these polymers have such a highhemical resistance that they do not dissolve in common sol-ents for membrane preparation. Therefore new synthetic routesf these polymers have focused mainly on increasing solubilitynd processability. Thermal and chemical stability of these poly-ers would be expected to meet the requirements of potential

eparation applications under harsh conditions such as hydrogenurification from steam reformer or carbon dioxide separation fromue gases (e.g., post-carbon capture process in coal-based powerlants). However, the gas permeation and separation data of suchigh-temperature polymer membranes are rarely reported mainlyue to difficulties in fabricating practical membrane geometriessing a conventional solution casting method because of lack of sol-bility in organic solvents. This fabrication challenge was recentlyvercome by our group [16] which adopted a post-fabricationolymer-modifying reaction [17–19] to obtain dense PBO and PBTembranes by the thermal arrangement of soluble aromatic poly-

mides containing ortho-positioned functional groups (e.g., –OHnd –SH). Membranes prepared with thermal rearrangement (TR)olymers showed excellent separation performance for CO2/CH4ixtures, with high selectivity and permeability due to an unusualicrostructure whose cavity size and distribution could be tuned

y the choice of template molecules and heat treatment. Moreover,R polymer membranes exhibited a high resistance to plasticizationn mixed-gas permeation experiments at 35 ◦C for carbon dioxideartial pressures as high as 20 atm [16].

In this study, we have further examined the gas separation per-ormance of TR polymer membranes by varying template moleculesnd heat treatment protocols. In addition, the mechanism of uniqueavity and distribution preferentially created during thermal rear-angement is analyzed from the standpoint of molecular dynamicsnd mechanics, together with experimental data such as positronnnihilation lifetime spectroscopy (PALS) and X-ray scattering data.

. Experimental

.1. Materials

Diamines: 2,2′-bis(3-amino-4-hydroxylphenyl) hexafluoro-ropane (bisAPAF) and 2,5-diamino-1,4-benzenedithiol (DABT)

ne Science 359 (2010) 11–24

were purchased from Tokyo Chemical Industry (TCI) Co., Ltd.(Tokyo, Japan).

Dianhydrides: 4,4′-(hexafuloroisopropylidene) diphthalicanhydride (6FDA), 3,3′,4,4′-bisphenyltetracarboxylic dianhydride(BPDA), 4,4′-oxydiphthalic anhydride (OPDA), 3,3′,4,4′-benzophen-onetetracarboxylic acid dianhydride (BTDA), 1,2,3,5-benzentetracarboxylic dianhydride (PMDA), and 1,4,5,8-naphthalenic tetracar-boxylic dianhydride (NTDA) were purchased from Tokyo ChemicalIndustry (TCI) Co., Ltd. (Tokyo, Japan) (PMDA, BPDA, OPDA andNTDA) and Aldrich (Milwaukee, WI, USA) (6FDA and BTDA).

Solvent: N-methyl-2-pyrrolidione (NMP, 99.0%) and m-cresol(99.0%) (for NTDA-based polyimide) was purchased from Aldrich(Milwaukee, WI, USA).

Sample codes: PIOFG-1 (6FDA + bisAPAF), TR-1 (thermallyrearranged polymer from PIOFG-1); PIOFG-2 (OPDA + bisAPAF),TR-2 (from PIOFG-2); PIOFG-3 (BTDA + bisAPAF), TR-3 (fromPIOFG-3); PIOFG-4 (BPDA + bisAPAF), TR-4 (from PIOFG-4);PIOFG-5 (PMDA + bisAPAF), TR-5 (from PIOFG-5); PIOFG-6(NTDA + bisAPAF), TR-6; PIOFG-S (6FDA + DABT), TR-S.

2.2. Synthesis of hydroxyl-containing polyimides

We prepared various polyimides containing ortho-positionedfunctional groups (PIOFGs), which can be synthesized fromdianhydrides and diamines in polar solvents (e.g., N-methyl-2-pyrrolidione (NMP)) [20]. The dianhydrides and diamines werepurified by vacuum sublimation or repeat crystallization prior topolymerization. NMP were stirred over CaH2 overnight and thendistilled under reduced pressure and stored over 4 Å molecularsieves under nitrogen. PIOFGs were prepared by a conventionaltwo-step process including formation of the hydroxyl-containingpoly(amic acids) followed by thermal imidization. For example, toa 250 mL flask purged with nitrogen was added bisAPAF (10 mmol)and NMP (25 mL). The mixture was stirred to form a clear, light yel-low solution at ambient temperature. To this solution of diaminewas added 6FDA (10 mmol) and NMP (25 mL) in order to achieve20% (w/v) solids content. The reaction was mechanically stirred for4 h at ambient temperature to give a viscous, light yellow poly(amicacid) solution. The poly(amic acid) solution was cast onto a glassplate. The cast films of the poly(amic acid) were dried under vac-uum at 60 ◦C, 100 ◦C, 150 ◦C, 200 ◦C, and 250 ◦C for 1 h, respectivelyand then thermally imidized to the hydroxyl-containing polyimideat 300 ◦C for 1 h under nitrogen. All other PIOFGs were preparedusing the same procedure.

2.3. Synthesis of TR polymer membranes

All polyimide membranes (4 cm × 4 cm in size; 20–30 �min thickness) were thermally treated at 300 ◦C, 350 ◦C, 400 ◦Cand 450 ◦C at a heating rate of 5 ◦C/min under argon flow(100 cm3[STP]/min) in a quartz tube furnace supported on an alu-mina holder plate [16,21]. The Ar flow was precisely controlled by amass flow controller (MKS Instruments, MASS-FLO, Andover, MA).A 4.5-cm-i.d., 70-cm-long quartz tube with a glass end cap wasused for the heat treatments in a muffle furnace equipped with fourheating elements to minimize the axial and radial temperature gra-dients. The length of the effective heating zone was 30 cm, and themaximum heating temperature was 1500 ◦C. Membrane were heldfor 1 h (or 3 h and 5 h) at the target temperature and then cooledslowly to room temperature.

2.4. Membrane characterizations

Attenuated total reflection Fourier transform infrared (ATR-FTIR) spectroscopy: The ATR-FTIR spectra of the PIOFGs and theTR polymers were measured using an infrared microspectrometer

Page 3: Thermally rearranged (TR) polymer membranes for CO2 separation

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IlluminatIR, SensIR Technologies, Danbury, CT, USA) in the rangef 4000–500 cm−1.

Thermal gravimetric analysis–mass spectroscopy (TGA–MS):GA was performed in flowing Ar by use of a TGA2050 ther-ogravimetric analyzer (TA Instruments, New Castle, DE, USA).

he amount of CO2 evolved during thermal rearrangement fromIOFGs to TR polymers was simultaneously detected by QMG 422igh-performance quardrupole mass spectrometer (Balzers Instru-ents, Liechtenstein). The PIOFG films were annealed at 100 ◦C forh in the TGA apparatus before each analysis to avoid the influencef humidity.

Wide-angle X-ray diffraction (WAXD): WAXD was measured tobserve the microstructural change of polyimides (PIOFGs) beforend after thermal rearrangement, using an X-ray diffractometerD/MAX-2500, Rigaku, Japan) which uses Cu K� radiation at a wave-ength 1.54 Å, operating in a 2� range of 5–60◦ with a scan rate of◦ min−1. From X-ray scattering data (using 2� of the broad peakaximum), the average d-spacing values of the PIOFGs and their

R polymers were calculated using Bragg’s law (d = �/2 sin �).Nitrogen adsorption/desorption isotherms: The apparent poros-

ty, total pore volume, average pore diameter and surface area ofR polymers were characterized from low temperature (77 K) nitro-en adsorption/desorption isotherms measured over a wide rangef relative pressure from 10−6 atm to 1 atm. Adsorption/desorptioneasurements were performed using a Micrometrics ASAP 2020

olumetric adsorption analyzer. High purity nitrogen (99.999%)as employed. Prior to the measurement, all samples wereegassed at 573 K for 4 h in a degas pot of the adsorptionnalyzer. The specific surface area of each sample was deter-ined from nitrogen adsorption isotherms by the standard

runauer–Emmett–Teller (BET) method and Langmuir method.icropore surface area and pore volume were calculated from

-plot and Horvath–Kawazoe equation [22]. The pore size distri-ution (PSD) was obtained from the regularization method basedn density functional theory (DFT) [23].

Gas permeation measurement: Molecular probe studies wereerformed using a constant-volume/variable-pressure method (soalled as a high-vacuum time lag method) at a feed pressure of60 Torr and a feed temperature of 25 ◦C [16] with different kineticiameters of gas molecules such as H2 (0.289 nm), CO2 (0.33 nm),2 (0.346 nm), N2 (0.364 nm) and CH4 (0.380 nm) [24]. Before theas permeation measurements, both the feed and the permeateides were thoroughly evacuated to below 10−5 Torr. The calibratedownstream volume was 30 cm3, the upstream and downstreamressures were measured using a Baratron transducer (MKS) withfull scale of 25,000 and 10 Torr, respectively. The pressure rise

ersus time transient of the permeate side equipped with a pres-ure transducer was recorded and passed to a desktop computerhrough an RS-232 cable. The linear slope of the pressure rise ver-us time provides the permeation rates of penetrating gases. Theermeability coefficient for a test gas is determined by multiply-

ng the permeation rate by the membrane thickness and can bexpressed by

= dp

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here P is the permeability represented in Barrer; dp/dt is the ratef the pressure rise under the steady state; V (cm3) is the down-tream volume; L (cm) is the membrane thickness; �p (cmHg) ishe pressure difference between the two sides; T (K) is the measure-

ent temperature; A (cm2) is the effective area of the membrane;

0 and T0 are the standard pressure and temperature, respectively.ll the gas permeation tests were performed more than five times,nd the standard deviation from the mean values of permeabil-ties was within ca. ±2%. Sample-to-sample reproducibility wasery high and within ca. ±5%. The effective membrane areas were

e Science 359 (2010) 11–24 13

4 cm2. The ideal separation factor (selectivity) for components 1and 2 is defined as the ratio of the pure-gas permeabilities of eachcomponent:

˛1/2 = P1

P2(2)

Positron annihilation lifetime spectroscopy (PALS) measure-ments: The PALS measurements were performed in nitrogen atambient temperature using an automated EG&G Ortec fast–fastcoincidence spectrometer [16]. The timing resolution of the sys-tem was 240 ps. The polymer films were stacked to a thicknessof 1 mm on either side of a 22Na–Ti foil source. There was nosource correction needed for the Ti foil (thickness 2.5 �m). Eachspectrum consisted of approximately 10 million integrated counts.The spectra were modeled as the sum of three decaying expo-nentials or as a continuous distribution. The shortest lifetime, �1,was fixed at 0.124 ns, which is characteristic of para-positroniumself-annihilation. The second lifetime, �2, was approximately0.35–0.45 ns for all samples, characteristic of free and trappedpositrons. The longer lifetime, �3, was >1 ns and attributed to anni-hilations of oPs in the free volume elements of the polymer.

Fraction free volume (Vf): The Vf is calculated using the followingequations [25–27]:

Vf = V − V0

V(3)

V = M

�(4)

V0 = 1.3 VW (5)

where V is the total molar volume of the repeating unit (cm3/mol),M the molar mass (g/mol) of the repeating unit and � the densityof the film (g cm−3), which was determined experimentally using atop-loading electronic Mettler Toledo balance coupled with a den-sity kit based on Archimedes’ principle. The samples were weighedin air and a known-density liquid, high purity water. The measure-ment was performed at room temperature by the buoyancy methodand the density was calculated as follows:

�polymer = w0

w0 − w1�liquid (6)

where w0 and w1 are the membrane weights in air and water,respectively. V0 is the volume occupied by the polymer chains(cm3/mol). V0 is assumed to be impermeable for diffusing gasmolecules. VW is the Van der Waals volume calculated using thegroup contribution method of Bondi [28].

3. Results and discussion

3.1. Thermal rearrangement mechanism

Generally the thermal rearrangements of hydroxyl-containingpolyimides to polybenzoxazoles (PBOs) are less well-known. A fewgroups reported that a series of hydroxyl-containing polyimides,including pendant hydroxyl groups ortho to the heterocyclicimide nitrogen, were synthesized by solution condensation ofvarious aromatic dianhydrides with bisaminophenols [16–19].Fig. 1 shows the proposed mechanism of thermal cyclizationfrom hydroxyl-containing polyimides to polybenzoxazoles. Themechanism includes a hydroxyl-containing polyimide with acarboxyl-benzoxazole intermediate. Hydroxyl-imide ring rear-ranged to a carboxyl-benzoxazole followed by decarboxylation,

in the range of 350–450 ◦C leads to a fully aromatic benzoxazoleproduct. It should be noted that the thermal rearrangement tem-perature depends on the type of template molecules for PIOFGsbecause the degree of thermal cyclization is closely related toinherent flexibility of PIOFG chains and restrictions imposed by
Page 4: Thermally rearranged (TR) polymer membranes for CO2 separation

14 H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24

F olyima

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ig. 1. Proposed mechanism of thermal rearrangement from hydroxyl-containing pcarboxyl-benzoxazole ring, and (c) a benzoxazole ring.

nterchain interactions. The intermediate state has bulky carboxyliccid groups leading to steric hindrance and chain disruption, so thatfter cyclocarboxylation some microvoids could be created with atatistical combination of meta- and para-linked species as well,ssociated with the original random chain conformations of PIOFGs,hat would result in more disordered chain structures. Intrinsi-ally, the thermally rearranged chains, polybenzoxazole chains cane supposed to be much more rigid than PIOFG chains, so thatny physical changes (e.g., the formation of microvoids due toegmental rearrangement and chain disruption) after the thermalonversion would be irreversible.

Fig. 2 shows the ATR-FTIR spectra of PIOFG-1 and TR-1 poly-ers (thermally treated at 450 ◦C, 1 h, argon atmosphere). Thermal

midization of PIOFG-1 was confirmed by the appearance of absorp-ion bands at 1788 cm−1 (b: symmetric C O stretching) and718 cm−1 (c: asymmetric C O stretching). There is also a strongnd broad absorption band between 3200 cm−1 and 3600 cm−1 (a:

Fig. 2. ATR-FTIR spectra of (a) PIOFG-1 and (b) TR-1 polymer films.

ide to polybenzoxazole: (a) a imide ring with ortho-positioned hydroxyl group, (b)

OH-phenyl). Thermal conversion from PIOFG-1 and TR-1 was con-firmed by absorption bands at 1474 cm−1 and 1059 cm−1 (d and e:typical bands of benzoxazoles [29]).

As mentioned above, the thermal rearrangement of hydroxyl-imide rings to benzoxazoles ring accompanies the evolution ofcarbon dioxide (2 mol CO2/PIOFG repeating unit). Most of PIOFGsprepared in this study was stable up to 350 ◦C, and the distinctweight loss occurred between 400 ◦C and 450 ◦C. At these tem-peratures, a polybenzoxazole structure formed due to a thermalcyclization reaction. Fig. 3 shows the weight loss of PIOFG-1 andthe corresponding mass (CO2) trace curve. Explicitly, the amountof CO2 evolution reached the maximum value at around 450 ◦C,suggesting that 450 ◦C is the effective temperature for thermal con-version of PIOFG-1 to TR-1. In our previous studies [16,20], theweight loss of five PIOFG samples (PIOFG-1–5) was experimentallymeasured using a TGA under Ar atmosphere and compared to the-oretical values. For PIOFG-1, the theoretical weight loss is 11.0%,which is very close to the experimental value (10.7%), meaning thatthe degree of thermal rearrangement is fairly high (PIOFG-2: 13.6%

(theoretical)/11.6% (experimental); PIOFG-3: 13.3%/10.9%; PIOFG-4: 13.9%/9.5%; PIOFG-5: 15.8%/13.7%). We also checked the thermalstability of TR polymer membrane in a thermo-oxidative state. Fig. 4shows the weight loss of TR-1 (thermally treated at 450 ◦C, 1 h) at

Fig. 3. The weight (%) of PIOFG-1 under N2 atmosphere and CO2 trace of massspectroscopy during the TGA experiment.

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H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24 15

Fig. 4. Thermal stability of TR-1 membrane under thermoxidative condition (airatmosphere).

Fig. 6. Simulated rotational distribution of dihedral angles in (a) a imide-phenylene withatures, A: 25 ◦C, B: 300 ◦C, C: 350 ◦C, D: 400 ◦C, E: 450 ◦C).

Fig. 5. Possible chain conformations of TR polymers derived from PIOFGs.

air atmosphere. TR-1 polymer exhibited excellent thermal stabilityup to 500 ◦C, suggesting that these TR polymers can be employedfor high-temperature gas separation applications.

3.2. Microvoids (free volume elements) formation

The solution-cast PIOFG membranes are completely amorphous(confirmed by X-ray scattering data), i.e., the polymer chains havea wide distribution of torsional angles and bond angles at a given

ortho-positioned hydroxyl group and (b) a benzoxazole-phenylene (A–E: temper-

Page 6: Thermally rearranged (TR) polymer membranes for CO2 separation

16 H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24

F polyms

tPmSfm(hefl

TP

ig. 7. Energy-minimized chain structure of simulated (a) PIOFG-1 and (b) TR-1ize = 1.1 Å).

emperature. It is assumed that the thermal rearrangement ofIOFG chains, in particular in the solid state, would result in the for-ation of microvoids caused by the chain disruption in stiff chains.

tatistically, as shown in Fig. 5, there are three types of chain con-ormations per a TR polymer chain, i.e., meta-meta, para–para, and

eta–para (=para–meta) catenated, supposing that the dihedral

torsional) angle between the imide ring and the ortho-positionedydroxyl-phenylene ring can be varied within a limited rotationalnergy barrier at a given temperature. Essentially, the rotationalexibility is significant with regards to the degree of thermal rear-

able 1hysical properties of various PIOFGs and TR polymers at 25 ◦C.

Sample Density (g cm−3) V (cm3 g−1) VW (cm3 g−1

PIOFG-1 1.536 0.665 0.430TR-1 1.293 0.773 0.439

PIOFG-2 1.453 0.688 0.459TR-2 1.271 0.787 0.473

PIOFG-3 1.469 0.681 0.455TR-3 1.304 0.767 0.469

PIOFG-4 1.482 0.675 0.457TR-4 1.240 0.806 0.470

PIOFG-5 1.478 0.677 0.443TR-5 1.362 0.734 0.457

ers, and free volume distribution of (c) PIOFG-1 and (d) TR-1 polymers (probe

rangement and chain packing efficiency. A single TR polymer chainwould have a random-sequenced chain conformation that is arbi-trarily mixed with the three different types of chain conformations.

Molecular dynamics and mechanics simulations provide morecomprehensive molecular visions of such a hypothesis. We built

atomic bulk models of the PIOFG-1 polymer with a single chainof 30 repeating units using information on the conformationalangle statistics and excluded volume effect. After several energy-minimization iterations and molecular dynamics steps [30], 99% ofthe experimental density was achieved at P = 0.1 MPa, indicating

) FFV Increment in FFV (%) d-Spacing (nm)

0.15965

0.5480.263 0.600

0.13464

0.5460.219 0.606

0.13157

0.5030.205 0.611

0.120102

0.5390.243 0.602

0.14828

0.5600.190 0.698

Page 7: Thermally rearranged (TR) polymer membranes for CO2 separation

H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24 17

F4

ttdcepnecract

ig. 8. Pictures of (a) PIOFG-1, (b) TR-1 treated at 400 ◦C, and (c) TR-1 treated at50 ◦C.

hat the simulated polymer chains may have well-relaxed struc-ures, similar to the actual system. Fig. 6 shows the actual rotationalistribution of the dihedral angle (C–N–Cph–Cph–OH, ϕ1 and ϕ2)alculated from the amorphous PIOFG-1 cells during canonicalnsemble (NVT) dynamics (t = 1500 ps) for different target tem-eratures in the range of 300–450 ◦C. The distribution peaks occurear to 45 ± 15◦ for ϕ1 (135 ± 15◦ for ϕ2), and the rotational barriernergy diminishes with increasing temperature in the simulation

ell, indicating that the majority of the OH-phenylene rings canotate more freely with increasing temperature. This denotes thatt more elevated temperatures, the three proposed conformationsan all occur depending upon the degree of rotation dominated byhe initial chain conformation of the precursor polymers. As such,

Fig. 9. Wide-angle X-ray scattering patterns of (a) PIOFG-1 and (b) TR-1 membranes.

the TR polymers can have either or both meta (m)- and para (p)-linked chain conformations. The quantitative amount of each chainconformation in a single chain would be different, depending uponthe original polymer chain.

Using the same amorphous cell construction procedure, we alsobuilt four types of amorphous polymer cells: the m–m, p–p, andm–p linked TR-1 series, and a single TR-1 chain of 30 repeatingunits randomly mixed with those three repeating units generatedat the same composition (each 10 repeating units) [16]. The molec-ular simulation study could not reflect the final microstructure bythe actual thermal rearrangement process, only but the effect ofthe physical properties on different chain conformations was esti-mated. The actual rotational distribution of the Cox–Cph bond (ϕ3)calculated from four amorphous TR-1 cells is shown in Fig. 6. Therange of torsional angles is 0 ± 10◦, revealing that the majority ofthe phenylene rings are coplanar with the oxazole units, irrespec-tive of the types of simulated amorphous TR-1 polymer cells. Thedensity of each TR-1 series amorphous cell was lower than that ofthe PIOFG-1 amorphous cell. Average values of simulated densitiesof at least 10 chains per polymer are 1.534 g cm−3 (for comparison,the experimental density is 1.536 g cm−3) for PIOFG-1, 1.342 g cm−3

for m–m linked TR-1, 1.341 g cm−3 for p–p linked TR-1, 1.351 g cm−3

for m–p linked TR-1, and 1.281 g cm−3 for mixed TR-1, respectively.The lowest density was obtained in mixed TR-1 chains, indicatingpoor chain packing efficiency. The experimentally determined den-sity of TR-1 polymer prepared at 450 ◦C is 1.293 g cm−3; this valueis closest that of the simulated mixed TR-1 chain. These simulationresults support our method to attain unexpected microvoid spacesin polymers in the solid state by spatial relocation attributable togeneration of dissimilar chain conformations.

Fig. 7 shows the simulated PIOFG-1 and TR-1 chains. AmorphousTR-1 chain (Fig. 7(b)) was obtained, based on the energy-minimizedchain structure of amorphous PIOFG-1 (Fig. 7(a)). The benzoxazolering in the TR-1 chain was regenerated from the most adjacenthydroxyl-phenylene rings and imide rings in the PIOFG-1 chain,

followed by iterative energy-minimization steps. As illustratedin Fig. 7(a) and (b), amorphous PIOFG-1 chain exhibited well-packed structure, but amorphous TR-1 chain derived from the samePIOFG-1 chain showed less-packed structure (more microporous
Page 8: Thermally rearranged (TR) polymer membranes for CO2 separation

18 H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24

F (c) PI

spapuoBptetsiama

3v

oaeeti

ig. 10. Positron annihilation lifetime spectroscopy of (a) PIOFG-1, (b) PIOFG-2 and

tructure), due to the formation of rigid, coplanar benzoxazole-henylene plane occurred in the limited-mobility-space. Fig. 7(c)nd (d) presents the accessible free volume elements of amorphousolymer cells simulated from PIOFG-1 and TR-1 polymer chains,sing a probe size of 1.1 Å. The accessible free volume elementsf TR-1 polymer are much larger than those of PIOFG-1 polymer.ased on these simulation approaches, chain disruption in stiffolymers, if occurring in a solid state, can give rise to microstruc-ural changes such as the evolution of microvoids, i.e., free volumelements. That is, modifications in the polymer chains that inhibithe polymer chain packing will tend to shift the distribution ofuch free volume elements. In addition, molecular modeling stud-es showed that there are different types of free volume in PIOFGnd TR polymers, and the gas transport mechanism in two polymerembranes will depend on the nature of the free volume elements

nd their size distribution.

.3. Change of physical properties (density, fractional freeolume)

In all cases, the densities of TR polymers were lower than thosef PIOFG polymers (see Table 1). It was reported that without

ny chemical reaction the density of aromatic polyimides in non-quilibrium state slightly increases owing to the thermal annealingffect—resulting in denser polymer chain packing as well as chargeransfer complex (CTC) [31]. Initially, we expected the minutencrease of density since total dimension of TR polymer films less-

OFG-3 at different heat treatment temperatures (300 ◦C, 350 ◦C, 400 ◦C and 450 ◦C).

ened within 3.0 ± 0.5% due to the release of CO2 molecules by thedecarboxylation, and the CTC phenomena witnessed by the colorchange from light yellow to dark purple as the target temperatureincreases (Fig. 8). Based on density data and by supposing that afull conversion would be achieved at 450 ◦C, we estimated the frac-tional free volume (Vf), a governing physical element liable for thetransport of small molecules in polymers, of PIOFG and TR polymersas listed in Table 1. For example, the calculated Vf values of PIOFG-1and TR-1 membranes are 0.159 and 0.263, respectively. The Vf val-ues of TR polymers are considerably larger as compared to thoseof polysulfone (0.159) and polycarbonate (0.166), and comparablewith those of high-free-volume glassy polymers such as poly(1-trimetylsilyl-1-propyne) (PTMSP) (0.29) and Teflon AF2400 (0.32).Undoubtedly, TR polymers contain a considerable amount of freevolume elements.

3.4. Change of average interchain distance (d-spacing)

Wide-angle X-ray scattering patterns of all PIOFGs and theircorresponding TR polymers exhibit that these polymers are com-pletely in an amorphous state without any crystalline phase. Asshown in Fig. 9, the 2� value at the maximum peak of a PIOFG film

are shifted to the left after thermal rearrangement at 450 ◦C, mean-ing that average interchain distances (d-spacing) of a PIOFG filmincreased after thermal rearrangement (e.g., PIOFG-1 (0.548 nm)and TR-1 (0.600 nm)). The d-spacing would not be taken as atrue measurement of the interlayer distance, but changes in the
Page 9: Thermally rearranged (TR) polymer membranes for CO2 separation

H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24 19

IOFG-

dabaat(p

3

crptPpcoaiipogo

Fig. 11. Nitrogen adsorption/desorption isotherms of (a) P

-spacing would serve as an indicator of the amount of roomvailable for penetrating small molecules through polymer mem-ranes. Not always, but usually the average interchain distance ofmorphous glassy polymers would be expected to decrease whennnealed at near or above glass transition temperature (Tg) owingo faster transition from non-equilibrium state to equilibrium statei.e., accelerated structural relaxation), resulting in increasing chainacking density [32].

.5. Positron annihilation lifetime spectroscopy (PALS)

PALS can be used to estimate the mean size and the con-entration of free volume elements occurred during the thermalearrangement. The free volume elements in amorphous glassyolymers are of crucial importance for elucidating gas transporthrough the membranes. PALS analysis of the PIOFGs (PIOFG-1,IOFG-2 and PIOFG-3) and the corresponding thermally rearrangedolymers shows that the TR polymers reconstituted from randomhain conformations undergo microstructural changes, dependingn the degree of thermal conversion at each designated temper-ture (Fig. 10). The oPs lifetime (ns) of all membrane couponsncreases as the temperature goes up from 300 ◦C to 450 ◦C. The oPs

ntensity (%) increases up to 400 ◦C, but decrease above this tem-erature. In PALS data, the longer oPs lifetime means the larger sizef the free volume elements. The larger oPs intensity indicates thereater concentration (or number) of free volume elements. ThePs intensity is related to the probability of oPs formation. Such

1, (b) TR-1 (400 ◦C), (c) TR-1 (450 ◦C) and (d) TR-1 (500 ◦C).

large changes are inconsistent with accepted belief that the chainpacking of glassy polymers typically increases due to the enhancedstructural relaxation at elevated temperature. An increase in meanpore size is accompanied by a decrease in number of free vol-ume elements—suggesting coalescence of two smaller cavities toform one larger cavity that can create interconnected free vol-ume elements. That is, the experimental data are an indicative ofthe existence of interconnected microcavities in TR polymer mem-branes.

3.6. Nitrogen adsorption/desorption isotherms

We measured the BET surface area and the single-point totalpore volume from nitrogen adsorption/desorption isotherms forPIOFG-1, TR-1-350, TR-1-450, and TR-1-500, as shown in Fig. 11.This measurement was usually applied for analyzing pore sizesand their distributions in porous inorganic materials rather thanorganic polymers. Indeed, dense polymer membranes are consid-ered to be non-porous in that free volume (i.e., cavities) elementsare not connected to each other, so this technique is seldomemployed to characterize the pore characteristics of polymericmembranes. There are a few studies about high-free-volume

glassy polymers (e.g., poly(phenylene oxide), poly(trimethylsilylpropyne) (PTMSP), and polymer with intrinsic microporosity (PIM))in the literature [33–35]. For instance, PTMSP possesses an openmicroporous structures due to a network of interconnected micro-cavities. The open porous structures also attribute to their highest
Page 10: Thermally rearranged (TR) polymer membranes for CO2 separation

20 H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24

polym

g[

nmTa(aiUaNstthtoBi

itemthtT

3

empc

P

ds

Fig. 12. (a) Gas permeability and (b) gas selectivity of TR-1

as permeation properties among all known polymeric materials36].

The sorption measurement of nitrogen shows a range fromon-porous polymers (PIOFG-1: BET surface area <20 m2 g−1) toicroporous polymers (TR-1-450: BET surface area >500 m2 g−1).

he nitrogen isotherms for TR-1-450 are irreversible and showhysteresis because of nitrogen desorption from the micropores

<2 nm), also indicating the existence of some larger microporesccessible only thorough the micropores. The BET surface areas 535 m2 g−1 and the single-point total volume is 0.296 cm3 g−1.sing the density of 1.293 g cm−3 for TR-1-450, we found that theccessible porosity is 40.9, i.e., the microvoid fraction is about 0.41.ote that the void fractions of the most open zeolites such as fauja-

ite, paulingite and the zeolite A families are 0.47–0.50 [24], so thathe current results are interesting. On the other hand, for TR-1-500,he nitrogen isotherms show a reversible type I isotherm and noysteresis, represented by a plateau that is nearly or quite horizon-al, and that may cut the relative pressure (P/P0) = 1 axis sharplyr may show a tail as the saturation pressure is approached. TheET surface area is 706 m2 g−1, the single-point total pore volume

s 0.343 cm3 g−1.The most likely cause of the hysteresis observed for TR polymers

s the activated diffusion of sorbate molecules into microcavitieshrough pre-existing throats of molecular dimensions. The differ-nce in equilibrium pressures for adsorption and desorption in theulti-throat cavity pore model becomes considerably larger than

hat for a single- or double-throat cavity, resulting in a very broadysteresis [37]. It seems that the multi-throat cavity model offershe most adequate representation of the free volume structure inR polymers.

.7. Gas permeation properties of TR polymer membranes

Gas transport through dense polymer membranes can belucidated from the solution-diffusion concept [2], in which per-eability coefficient (PA) of a gas (cm3(STP)cm/cm2 s cmHg) is the

roduct of diffusion coefficient (DA) of a gas (cm2/s) and a solubilityoefficient (SA) of a gas (cm3-gas/cm3-polymer cmHg).

A = DASA (7)

Ideal selectivity (˛A/B) of gas A to gas B is both due to theifferences in diffusion coefficient (diffusivity selectivity) and inolubility coefficient (solubility selectivity).

er membrane as a function of heat treatment temperature.

Our previous study [16] shows that the rearrangement of PIOFGchains was thermally triggered at around 350 ◦C, and full con-version to TR polymers occurred at around 450 ◦C, estimated byTGA–MS data. When the PIOFGs were thermally treated below450 ◦C, the TR membranes were considered to contain bothunchanged segments (e.g., polyimide) and thermally rearrangedsegments (e.g., polybenzoxazole). Therefore, the gas permeationproperties depend strongly on the final heat treatment tem-perature. As shown in Fig. 12, the gas permeability increasesdramatically and the gas selectivity decreases by increasing theheat treatment temperature. That is, the degree of thermalrearrangement of the precursor PIOFGs appreciably affects the gas-accessible free-volume elements of the TR polymer membranes. ForTR polymers treated at 450 ◦C, the order of permeability is as fol-lows: CO2 > H2 > He > O2 > N2 > CH4 and the gas permeability valuesare very high. Usually polyimide membranes show higher H2 per-meability than CO2 permeability, which is consistent with the orderof increasing kinetic diameters of the penetrant molecules, show-ing that diffusion selectivity plays an important role for polyimides.Current TR precursor polymers, PIOFG membranes follow this orderof permeability. By increasing heat treatment temperature of PIOFGmembranes, the differences between H2 and CO2 permeabilitiesare gradually reduced in the range up to 400 ◦C. When fully con-verted TR polymer is obtained at 450 ◦C, the order of permeabilityis changed as shown in Fig. 11(a). This unusual permeation behav-ior is due to the enormous fractional free volume of TR polymers,suggesting that much higher CO2 solubility than H2 solubility in TRpolymers contributes to higher CO2 permeability through the TRpolymer membranes.

The gas selectivities (e.g., CO2/CH4, CO2/N2, O2/N2 and N2/CH4)decreases with increasing the heat treatment temperatures. AllPIOFG and TR polymer membranes exhibit high CO2/CH4 andCO2/N2 separation properties. Although the selectivities decreasefrom PIOFG to TR polymer, TR polymer membranes will be promis-ing membrane materials for CO2 separation in that they have muchhigher CO2 permeability with still high CO2/CH4 and CO2/N2 selec-tivities than conventional glassy polymers (e.g., cellulose acetateand polyimides) used for CO2 separation. According to our previ-ous study [16], much more importantly in TR polymer membranes

for CO2 separation is the fact that TR polymer membrane, asfully converted, did not show a significant CO2/CH4 selectivityreduction largely due to CO2 plasticization usually occurred inthe conventional glassy polymers. We believe that this might beclosely related to the origin of PIOFG membranes in preparation
Page 11: Thermally rearranged (TR) polymer membranes for CO2 separation

H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24 21

Fm

aRftobo6pbmbfiectn

cbtmsfrdtfitpciCi

indicating that such a temperature range may not change consid-erably the microstructures of thermally stable polyimide affectingthe gas diffusion and solubility, if not any large physical and spatialchanges in polymer chains as in TR polymers.

ig. 13. CO2 permeability and CO2/CH4 selectivity of fluorinated polybenzoxazoleembranes as a function of degree of thermal cyclization (ref. [39]).

nd thermally activated crosslinking during the heat treatment.ecently we have found that TR polymer membranes show dif-

erent gas permeability and gas selectivity depending on theype of imidization of PIOFGs. It is known that polyimide filmsbtained by the thermal imidization of poly(amic acid) mem-ranes exhibited lower permeability than polyimide membranesbtained by chemical imidization. According to the literature [38],FDA-pODA membranes prepared by chemical imidization hadermeability of PCO2 = 17 Barrer, and selectivity of CO2/CH4 = 62,ut membranes prepared by thermal imidization had lower per-eability and selectivity: PCO2 = 6 Barrer, CO2/CH4 = 43. This might

e due to the fact that heating at temperatures above 350 ◦Cor extended periods of time may result in crosslinking. Mostmides become insoluble when exposed to these conditions. How-ver, it is not yet clear whether this insolubility is caused byrosslinking or to enhanced chain packing. Crosslinking at elevatedemperatures would most likely proceed by a free radical mecha-ism.

The increase in gas permeability during different thermalonversion processes was also found in the polybenzoxazole mem-ranes from poly(o-hydroxyamide)s. Okamoto et al. [39] reportedhe gas permeation properties of fluorinated polybenzoxazole

embranes prepared by casting N,N-dimethylacetamide (DMAc)olutions of the precursor poly(o-hydroxyamide)s on glass platesollowed by the thermal cyclohydration. In this case, condensationeaction occurs followed by the evolution of water molecules. Usingifferent curing temperatures, various degrees of thermal cycliza-ion from o-hydroxyamide to benzoxazole ring was achieved innal polymer structure, and then measured various gas permeabili-ies and selectivities. Based on their data [39], we re-plotted the CO2

ermeability and CO2/CH4 permselectivity as a function of thermalyclization degree in Fig. 13. The CO2 permeability increased byncreasing the degree of thermal cyclization from 0% to 100% whileO2/CH4 selectivity decreased slightly as the curing temperature

ncreased. This trend is similar to our case, implying that chain rear-

Fig. 14. CO2 permeability and CO2/N2 selectivity of a commercial polyimide(Matrimid) membrane treated at different temperatures (ref. [31]).

rangements of polymers in a glassy state can lead to a significantchange in free volume elements determining the gas transport andseparation.

In contrast, Barsema et al. [31] reported the gas permeationproperties of commercial polyimide membranes thermally treatedat elevated temperatures. They observed the change of various gaspermeabilities (e.g., O2, N2 and CO2) through thermally treatedMatrimid membranes at different temperatures in the range of300–525 ◦C (Fig. 14). The polyimide membranes annealed below475 ◦C did not show any remarkable change in gas permeability,

Fig. 15. Effect of thermally rearrangable block on O2 permeability and O2/N2 selec-tivity of copolymer membranes.

Page 12: Thermally rearranged (TR) polymer membranes for CO2 separation

22 H.B. Park et al. / Journal of Membrane Science 359 (2010) 11–24

Table 2Gas permeation data of TR-1 membranes doped with various dopants at 25 ◦C.

Sample CO2 H2 He O2 N2 CH4

TR-1-HCl 912 1231 816 237 41 14.8TR-1-HCl (dedoping) 759 1062 696 250 56 20.4

achu[aswmm

mobpHC0eswibohbara

Fm

TR-1-HCl (redoping) 702 942 639 203 44 17.0TR-1-HNO3 325 470 314 105 57 54.8TR-1-H3PO4 295 738 559 82 12 3.3TR-1-HPF6 5 30 53 2 0.28 1.3

To further see the effect of chain rearrangement on the gas sep-ration performance, we considered the copolymers consisting ofommon polyimide and PIOFG. For this purpose, we designed oneomopolymer and three copolymers using two different buildingnits, [A] and [B] (mole ratio of [B]/[A] = 0/10, 1/9, 5/5 and 9/1). UnitA] is BTDA-ODA (as a polyimide) and block [B] is BTDA-bisAPAF (asPIFOG). All the copolymers were thermally treated at 450 ◦C. As

hown in Fig. 15, the O2 permeability increased as more of block [B]as added, which decisively indicated that indeed these copoly-ers originating from building units like [B] prove to be highlyicroporous polymers like TR polymers from a homopolymer.To exploit the microporous characteristics of current polymer

embranes, we carried out the acidic doping by changing the sizef the dopant species because current family of polymers containsasic site (=N–) in benzoxazole and benzothiazole rings [16]. TR-1olymer membranes were each doped with HCl, HNO3, H3PO4 andPF6 in 10 M aqueous solution (the order of thermochemical radii:l− = 0.168 nm < NO3

− = 0.200 nm < H2PO4− = 0.213 nm < PF6

− =.242 nm). During the doping process, we observed some inter-sting features. When HCl was used as a dopant (10 M aqueousolution), the doping level reaches at about 108 mol% within 3 h,hich is higher than 54 mol% of polybenzimidazole (PBI) contain-

ng two basic sites per repeating unit. Basically the reactivity ofenzoxazole to acid molecules is known to be lower than thatf benzimidazole since the oxygen atom (3.5) of benzoxazole

as higher electronegativity than the nitrogen atom (3.0) ofenzimidazole ring so that higher electronegativity of oxygentom prohibits the delocalization of electrons in benzoxazoleing and compensates the electron-donating effect of oxygentom with lone pair electrons. In other words, the doping level

ig. 16. Relationship between O2 permeability and O2/N2 selectivity of TR polymerembranes tested in this study with upper bounds.

Fig. 17. Relationship between CO2 permeability and CO2/CH4 selectivity of TR poly-mer membranes tested in this study with upper bounds.

of benzoxazole to acid molecules was expected to be relativelylower than that of benzimidazole with higher basic nature. Coun-terintuitively, our observed high doping level provides furtherevidence of microporous structures in these TR polymer mem-branes imparting a large surface area per volume. As listed inTable 2, the gas permeability and gas selectivity in TR polymermembranes show significant changes strongly depending on thetypes of dopants, indicating that the gas separation performanceof TR polymer membranes can be further tuned by doping of smallmolecules.

Figs. 16–20 show the relationship between gas permeability and

gas selectivity in TR polymer membranes studied in our group bycomparing the data with Robeson’s plots in 1991 and 2008 [40,41].Most of TR polymer membranes showed outstanding gas separa-tion performance in all interesting gas pairs shown in these figures.

Fig. 18. Relationship between CO2 permeability and CO2/N2 selectivity of TR poly-mer membranes tested in this study with upper bounds.

Page 13: Thermally rearranged (TR) polymer membranes for CO2 separation

H.B. Park et al. / Journal of Membran

Fm

GbCtmmrThueautt

Fm

ig. 19. Relationship between H2 permeability and H2/CO2 selectivity of TR polymerembranes tested in this study with upper bounds.

enerally the gas separation performances surpass the 1991 upperound of polymeric gas separation membranes. In particular, theO2/CH4 plot are worth noting as most of TR polymer membranes inhis study outperform even the new 2008 upper bound of polymeric

embranes, indicating that these TR polymers could be prospectiveembrane materials for CO2 separation applications such as natu-

al gas processing, landfill gas recovery, and enhanced oil recovery.he CO2/N2 separation of TR polymer membranes also show theigh selectivity with outstanding CO2 permeability, which can besed for CO2 capture from flue gas in post-combustion process withxcellent thermal stability. A noteworthy feature is the N2/CH4 sep-

ration of TR polymer membranes. Generally the N2/CH4 separationsing polymeric membranes is believed to be difficult because ofhe molecular properties of the N2 and CH4 and the nature ofhe gas permeation process [42]. Nitrogen molecules are slightly

ig. 20. Relationship between N2 permeability and N2/CH4 selectivity of TR polymerembranes tested in this study with upper bounds.

e Science 359 (2010) 11–24 23

smaller than those of methane, and thus diffusion favors nitrogen.However, methane is more condensable than nitrogen, and so sol-ubility favors methane. As such, the solubility of N2 is lower thanthat of CH4 due to lower critical temperature of N2, and the sol-ubility selectivity of N2 over CH4 is within the range of 0.2–0.4[42]. On the other hand, the diffusivity of N2 is higher than thatof CH4 due to the smaller kinetic size of N2. Based on solution-diffusion mechanism, consequently, N2/CH4 selectivities tend tobe low in polymer membranes. Among rubbery and glassy poly-mers, the diffusivity selectivity of N2/CH4 is small and tends to beovercome by the solubility selectivity. Consequently the perme-ability of N2 is smaller than that of CH4 for almost all polymersthat have been reported. Generally, polyimides are unique poly-mer systems showing reverse permselectivity of N2/CH4 becausethe diffusivity selectivity is much higher than in other polymerswhile the solubility selectivity is similar to that in other materials.This trend continues the N2/CH4 separation of TR polymer mem-branes despite their high gas permeability. The strong size andshape selective function of the TR polymers is believed to be dueto the limited mobility of TR polymer backbone segments, whichleads to a relatively narrow distribution of free volume elementssizes responsible for selective diffusion.

4. Conclusions

Current unexpected physical phenomena in TR polymers are ofgreat significance in that the random chain conformations occurredin condensed polymeric phase can result in tuned microvoidswhich contribute to performance enhancement in the selectivemolecular transport through the polymer membranes. The great-est advantage of the current methodology is to readily controlaverage interchain spacing and free volume elements that directlylead to molecular sieving effects, using different thermal proto-cols and template molecules. This enables us to make membranessuitable for specific gas separation applications. The gas perme-abilities of almost all TR polymer membranes are enhanced bytwo orders of magnitude over those of the original polymers andtypical glassy polymers, confirming the presence of larger inter-connected free volume elements. The gas permeabilities are lowerthan those of PTMSP membrane, the highest gas permeable poly-mer, but gas selectivities toward important gas separations (e.g.,CO2/CH4, CO2/N2) are 200–300% higher than in PTMSP and alsocomparable to or slightly lower than in inorganic molecular sievemembranes. These results support the claim that these TR poly-mer membranes posses “just right” microporous characteristicsbetween extreme high-free-volume glassy polymers and molec-ular sieve inorganic materials. The free volume of TR polymersis believed to be a three-dimensional network of intermolecu-lar microcavities. These are accessible for small gas moleculesdown to cryogenic temperatures, the interconnected microcavi-ties are analogous to micropores in certain solid adsorbents (e.g.,carbon molecular sieves) with respect to gas sorption and diffu-sion. The gas transport in TR polymers occurs through these newlycreated micropores. It is believed to be responsible for the molec-ular sieving which is common for separation of gaseous mixturesby glassy polymer membranes, the narrowest part of microporesplaying the role of a molecular size caliber and the chemical com-position of the macromolecules that form the pore walls. Thispeculiarity of a free volume structure may account for both theoutstanding permeability of TR polymers with fast diffusion of

gases and their still high permselectivity in separations of smallmolecules. In summary, we have developed a new method tocreate tailored microvoids with well-connected morphology forenhanced molecular diffusivity and sorptivity in amorphous poly-mer materials. This discovery should open the door for further
Page 14: Thermally rearranged (TR) polymer membranes for CO2 separation

2 mbra

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[meric membranes, J. Membr. Sci. 62 (1991) 165–185.

[41] L.M. Robeson, The upper bound revisited, J. Membr. Sci. 320 (2008) 390–400.[42] T.C. Merkel, I. Pinnau, R. Prabhakar, B.D. Freeman, Gas and vapor transport prop-

4 H.B. Park et al. / Journal of Me

reakthroughs based on the significance of void size, shape, andnterconnectivity in amorphous materials for adsorption, storagend separation.

cknowledgements

This work was financially supported by the Carbon Dioxideeduction and Sequestration Center, one of the 21st Centuryrontier R&D Programs funded by the Ministry of Science and Tech-ology in Korea.

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