Plasticization Suppression in Grafted Polyimide−Epoxy Network

Ind. Eng. Chem. Res. 2007, 46, 8183-8192

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Plasticization Suppression in Grafted Polyimide-Epoxy Network Membranes Xavier J. Duthie,† Sandra E. Kentish,*,† Clem E. Powell,† Greg G. Qiao,† Kazukiyo Nagai,†,‡ and Geoff W. Stevens† CooperatiVe Research Centre for Greenhouse Gas Technologies (CO2CRC), Department of Chemical and Biomolecular Engineering, The UniVersity of Melbourne, ParkVille, 3001, Victoria, Australia, and Department of Applied Chemistry, Meiji UniVersity, Kawasaki, 214-8571, Japan

This paper considers the blending of epoxide and diamine components into a glassy polyimide structure (2,2′-bis(3,4′-dicarboxyphenyl) hexafluoropropane dianhydride-2,3,5,6-tetramethyl-1,4-phenylenediamine (6FDATMPDA) intended for use as a gas separation membrane. It is shown that the diamine component can react both with the epoxide component to form an epoxy network and with the polyimide backbone itself, leading to a complex grafted structure. Bulk density and X-ray diffraction results show that this leads to a denser, more amorphous membrane structure. This effect appears to be independent of the diamine structure. Gas permeability is reduced and selectivity increases, consistent with established free volume theory. Importantly, however, plasticization resistance also increases. This implies that membrane performance will be more robust when exposed to condensable gases such as carbon dioxide. In this particular case, the increase in plasticization resistance is possibly insufficient to warrant the loss in permeability, particularly when compared to the relatively high permeability measured for diamine-cross-linked polyimide. Introduction A significant limitation of glassy polymer separation membranes is the phenomenon known as plasticization.1,2 The diffusion of condensable gases and vapors into the membrane structure causes a swelling effect which increases the polymer’s fractional free volume at high partial pressures of the penetrating gas. In turn, this results in an overall increase in membrane permeability and a loss in selectivity and leads to time-dependent membrane performance. In their dialogue on the future of gas separation membranes, Koros and Mahajan3 note that cross-linking of polymer membranes is one method that can provide plasticization resistance. Indeed, in 1999 Staudt-Bickel and Koros4 investigated the suppression of CO2-induced swelling and plasticization of glassy polyimide membranes for CO2/CH4 separation by chemical cross-linking. They utilized 6FDA-based polyimides and copolyimides with pendent carboxylic acid groups, which were blended with ethylene glycol, before inducing the cross-linking reaction thermally. The cross-linking reagent was included in the polymer via two mechanisms: either by blending with the polymer solution prior to casting or via absorption into the cast polymer film in a methanol/ethylene glycol solution. It was found that this method of cross-linking was successful at suppressing plasticization and swelling, and the actual CO2/ CH4 selectivity was found to increase.4 Most significantly, however, it was found that the CO2 permeability did not decrease by any significant amount with cross-linking. It was proposed that this was due to the insertion of a -(CH2)2- group as the cross-linker, which offset the reduced chain mobility by increasing the overall free volume of the system. No investigation, however, was made of the extent of the plasticization resistance over extended time frames. The patent of Hayes5 illustrated in 1991 that polyimides may be cross-linked with diamine compounds. Improvements * To whom correspondence should be addressed. E-mail: sandraek@ unimelb.edu.au. † The University of Melbourne. ‡ Meiji University.

were reported in oxygen/nitrogen selectivity. Chung and co-workers6-8 have also investigated the cross-linking of polyimides, with methanol/diamine solutions. Examples of the diamine compounds utilized include both linear7 and aromatic amines5,8 with suppression of plasticization exhibited in all cases. However, in this work the membrane permeability is also substantially reduced by the crosslinking procedure, with a concurrent increase in membrane selectivity. A cross-linked epoxy resin can similarly be formed from the addition of a diamine compound to an epoxide. It has been shown previously that the addition of such an epoxy resin can enhance the physical and/or chemical interactions in polymer materials.9 In particular, the blending of a polyimide with an epoxy resin in a dense membrane formulation appeared to suppress plasticization with elimination of hysteretic behavior in CO2 permeability isotherms.10 It is the objective of the present paper to further examine the addition of a range of cross-linkable epoxy compounds to the well-known polyimide system 2,2′bis(3,4′-dicarboxyphenyl) hexafluoropropane dianhydride-2,3,5,6tetramethyl-1,4-phenylenediamine (6FDA-TMPDA) (Figure 1) to form a cross-linked membrane structure. Background and Theory. The reaction of an epoxide with a diamine to form an epoxy polymer network is generally considered to proceed by three related reactions,11-13 each of which involves a nucleophilic ring opening of the oxirane group. These reactions are summarized in Figure 2. The “curing” of epoxy/amine systems often involves heating through a stepwise temperature profile. Initial curing begins at a temperature near the cure temperature, TC, which allows the two reactions detailed in Figures 2a and 2b to proceed. For steric reasons, reaction of the primary amine occurs faster than that of the secondary amine.12 Following this initial cure, a postcure process is initiated at a temperature greater than the TC.12 Primarily, the postcure process is designed to allow the unreacted secondary amines to be converted. The etherification reaction (Figure 2c) proceeds slowly compared to the amine-based conversion.12 Indeed, some studies have suggested that this reaction occurs only in epoxy-excess

10.1021/ie070689h CCC: $37.00 Published 2007 by the American Chemical Society Published on Web 10/31/2007

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Figure 1. 2,2′-Bis(3,4′-dicarboxyphenyl) hexafluoropropane dianhydride-2,3,5,6-tetramethyl-1,4-phenylenediamine (6FDA-TMPDA) polyimide.

Figure 2. Reactions in an epoxide-amine system.

systems and proceeds after the primary and secondary amines have been consumed.11-13 In a polyimide/epoxy system it is also possible for the amine reagent to react directly with the polyimide chain as described recently in the literature.5-7,14 The imide ring is opened to form two amine groups (Figure 3). If the reaction occurs between a polyimide and diamine, it is possible that this ring opening will indeed cross-link neighboring polyimide chains. Some workers have observed that this reaction is reversed at temperatures comparable to the epoxy curing temperature TC, described above.6,7 Our recent work, however, pointed out that, during the reverse reaction, there is no theoretical reason why only the original diamine should re-form.14 Using a monoamine and 6FDA-TMPDA, we illustrated that it is also possible for the imide ring closure to occur at the carbonyl group, either in the meta or para position relative to the hexafluoroisopropylidene moiety. Reaction at the meta linkage was shown to cleave the polyimide main chain, as evidenced by the subsequent reduction of molecular weight (Figure 3d). Alternatively, ring closure at the para position was found to restore the diamine moiety, Figure 3c, with no change in polyimide molecular weight recorded. It is evident from Figures 2 and 3 that in an epoxy/polyimide system a range of structures may result. Figure 4 illustrates a possible reaction sequence between a polyimide and epoxy system during formation and curing. The epoxide in this case is tetraglycidyldiaminodiphenylmethane (TGDDM). During the

formation period at ambient temperature, the diamine is capable of reaction with the carbonyl groups on the imide ring, yielding a tetraamide compound, Figures 4a and 4b. Upon heating, this reaction will be partially reversed, resulting in a mixture of diamine, TGDDM, polyimide, partially cross-linked polyimide, polyimide with pendent primary amine groups, and polyimide with terminal amine groups (Figures 4b and 4c). Simultaneously, the oxirane rings in the epoxide will begin to react with any primary amines and, later, secondary amines. Since there are large numbers of primary amines, and that each molecule of TGDDM possesses four oxirane functional groups, a very large number of potential structures can be generated. Figure 4d illustrates this by displaying a TGDDM which has reacted with two diamine molecules, one polyimide with a pendent amine group and one polyimide with a terminal amine group. In reality, the complexity will be much greater, as surplus oxirane groups will react with the secondary amines, leading to a much larger unit. For these reasons, it is proposed that the membranes examined in this work are representative of polyimide structures randomly grafted into an epoxy network via diamine crosslinking. Experimental Section Membrane Casting. 6FDA-TMPDA (PI) was synthesized by reaction between 4,4′-(hexafluoroisopropylidene)diphthalic

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Figure 3. Reactions of a polyimide, (a), with an amine to open the imide ring, (b), as well as possible ring-closing reactions, (c) and (d), that could occur upon heating.

anhydride (6FDA, Alfa Aesar, purity > 96.0%) and 2,3,5,6tetramethyl-1,4-phenylenediamine (TMPDA, TCI, purity > 98.0%)) in anhydrous N-methylpyrolidone (AR grade) to give a polyamic acid, which was subsequently imidized in situ with triethylamine and acetic anhydride to yield the desired polymer. The epoxide compound used was the same compound used by Nagai et al.10 in their initial investigation of polyimideepoxysystems: Tetraglycidyldiaminodiphenylmethane(TGDDM)s This is the epoxide used by commercial supplier Huntsman in its Araldite series of products. The structure of TGDDM is given in Figure 5. TGDDM was used as received from Aldrich. The Huntsman Araldite formulation uses diaminodiphenylsulphone as the diamine component. However, in the present work, the following diamines were considered as these dissolved more readily in the casting solvent, dichloromethane (Figure 5): 2,3,5,6-tetramethyl-1,4-phenylenediamine, TMPDA (TCI >

98.0% purity); 1,3-phenylenediamine, PDA (Fluka >99.0% purity); 4,4′-(9-fluorenylidene)dianiline, FDA (Sigma Aldrich >99.0% purity). These three diamines were selected to provide a range of packing abilities. The smaller PDA would be expected to interpenetrate the epoxy network easily, leading to a dense lowpermeability structure. Conversely, the bulkier diamines might lead to a more open polymer network. It was hoped that the resulting high free fractional volume would allow the polyimide permeability to be retained in spite of the interpenetration of the epoxy network. Membranes were prepared by solvent casting using a premixed system; epoxide, polyimide, and diamine were all separately dissolved in dichloromethane at 2.5 g/100 mL. Prior to casting, the epoxide solution (TGDDM dissolved in dichloromethane) was combined with the 6FDA-TMPDA polyimide

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Figure 4. Possible mechanisms for cross-linking of 6FDA-TMPDA via epoxide-diamine addition.

Figure 5. Structures of the epoxy components used in this study.

solution in a volumetric flask. Subsequently, the appropriate diamine solution was added and the entire mixture agitated manually. The order of mixing was followed scrupulously, as

alternative methods may lead to inhomogeous membranes. The composite solution was then solution-cast using glass casting rings on glass plates, which were covered with Petri dishes,

Ind. Eng. Chem. Res., Vol. 46, No. 24, 2007 8187 Table 1. Epoxide and Diamine Concentrations in Polyimide/Epoxy Films

epoxide (wt %)

diamine (wt %)

total epoxy precursora (wt %)

0.60 1.2 4.8 9.6 15.9 22.3 28.7

0.15 0.3 1.2 2.5 4.1 5.7 7.3

0.75 1.5 6.0 12 20 28 36

0.008 0.016 0.069 0.147 0.30 0.50 0.75

TMPDA

0.54 1.08 4.3 8.6

0.21 0.42 1.7 3.4

0.75 1.5 6.0 12

0.007 0.015 0.062 0.133

FDA

0.41 0.82 3.3 6.6

0.34 0.68 2.7 5.4

0.75 1.5 6.0 12.0

0.006 0.011 0.047 0.101

diamine type PDA

mole ratio diamine:PI repeat unit (mol/molPI_Repeat_Unit)

a Total epoxy precursor includes both epoxide and diamine compounds in a 1:1 mole ratio.

providing for slow evaporation at room temperature over approximately 24 h. Final membrane thicknesses ranged from from 40 to 60 µm. A range of epoxy concentrations were considered (Table 1). The ratio of TGDDM epoxide to diamine was equimolar in all cases, with these materials grouped under the term “epoxy precursor”. Some membranes were also prepared from polyimide/diamine mixtures (i.e., no epoxide) to allow a comparison to the polyimide/diamine systems already reported in the literature. Membrane Curing. The cast films were cured at ambient pressure according to the following protocol: (i) 30 min at 80 °C; (ii) 30 min at 100 °C; (iii) 90 min at 120 °C; (iv) 120 min at 200 °C. Heating rates between each step were at 10 °C min-1. This protocol matches that recommended by Huntsman (Araldite MY 720 Data Sheet) except that the final postcure temperature was increased from 177 to 200 °C to ensure extensive reaction of the two components. Following film curing, all membrane samples were dried at 150 °C and 0.1 mmHg for 24 h to ensure complete removal of the solvent. Analysis. Membrane densities were determined using calcium nitrate solutions and a hydrometer, as described in the ASTM D 1505-96.15 Estimates of the occupied volume for the polyimide-epoxy mixtures were made by determining this parameter separately for each component using established methods.16 This data was used in the estimation of the fractional free volume (FFV).1,17 Gel fractions were estimated by immersing a known mass of film sample in dichloromethane for 24 h, after which the insoluble material was removed and washed again in dichloromethane. Following vacuum drying at 80 °C overnight, the mass of insoluble material was measured. The permeabilities of the polyimide/epoxy membranes were determined using a constant volume, variable pressure gas permeation apparatus operating at 35 °C, as described previously. 18 Degradation temperatures and rates were determined on a Perkin-Elmer Diamond Thermogravimetric Analyzer, over the temperature range 50-550 °C, at a heating rate of 10 °C min-1 under nitrogen. Glass transition temperatures were determined using differential scanning calorimetry using a Perkin-Elmer Diamond DSC. Samples of approximately 5 mg were heated at a rate of 20 °C min-1, also under nitrogen, with the Tg determined during the first temperature scan. Polymer chain

spacings were estimated using a Ringuku 40 kV wide-angle X-ray diffractometer. With use of a Cu KR source (λ ) 1.54 Å), samples were analyzed over a range of 5-50°, with an increment size of 0.02° and scan rate of 2° min-1. Results and Discussion IR Spectra. IR spectroscopy was used to confirm reaction of the epoxide with diamine and subsequent formation of epoxy material. Initially, FTIR-ATR spectra of uncured TGDDM-PDA as well as cured TGDDM-PDA epoxy samples were studied in the absence of any polyimide (Figure 6). A reference peak at 1515 cm-1 was utilized, which was used by Musto et al.19 in their spectroscopic study of epoxy formation. It is attributed to a ring semicircle stretching mode of the diphenylmethane group within the TGDDM molecule.19 The literature reports the asymmetric deformation of the oxirane group for TGDDM as between 91619 and 910 cm-1.20 However, Nyquist and Fiedler illustrate that the precise frequency varies according to the solvent composition.21 In Figure 6, this peak occurs in the uncured TGDDM-PDA sample at 906 cm-1. The loss of this peak upon curing indicates the destruction of the oxirane group.19,22 Further evidence of epoxy formation is found at 1050 cm-1, where, similar to Musto et al.,19 significant absorbance increases are observed. Other workers have noted that monosubstituted oxirane groups such as those present in TGDDM also absorb around 800 cm-1.23 We observed significant depression in the absorbance at around 800 cm-1 upon curing, which parallels findings of Wu and Soucek.24 IR spectra obtained for a PI/TGDDM-PDA 0.745 mol/mol film, before and after curing, were then studied. The absorption peak at 1240 cm-1 (C-F group) was used as an internal standard for polyimide-containing samples, consistent with the work of Shao et al.7 As shown in Figure 7, the reduction of the peak at 906 cm-1 in the epoxy system is consistent with reaction of the epoxide. This peak, however, is slightly obscured by part of the polyimide spectrum, which has been noted previously.25 Infrared absorption at 800 cm-1 also decreases notably after curing, confirming TGDDM conversion. Finally, in results paralleling those obtained for the neat system, notable absorbance increases were also found around 1050 cm-1, further confirming epoxy formation. Macrostructural Analysis. Addition of increasing concentrations of epoxy compounds to the base polyimide formulation led to more intensely colored films and higher insoluble gel fractions (Figure 8). Indeed, addition of 20 wt % TGDDMPDA epoxy (0.3 mol of diamine/mol of polyimide repeat unit) led to over 70% of the film mass being insoluble. It is unlikely that physical trapping of the polyimide within the epoxy network could cause such an effect, suggesting that not only did epoxy material form but was indeed grafted into the polyimide macrostructure. Films cast with TGDDM-TMPDA and TGDDMFDA epoxy systems gave comparable gel fractions. Use of the largest diamine (FDA) resulted in slightly lower gel fractions at each wt % epoxy level. Conversely, polyimide films cast with identical PDA concentrations as used in the PI/TGDDM-PDA 36 wt % film but without the TGDDM epoxide were found to have no measurable gel content. Bulk density results showed similar trends. The films formed using the larger FDA diamine resulted in the smallest increase in density at constant wt % epoxy. Figure 9a also shows that the polymer formed from FDA results in the highest fractional free volume at a given wt % epoxy level. However, when this data is replotted on the basis of the number of diamine molecules available per polymer repeat unit (Figure 9b), it is clear that

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Figure 6. FTIR-ATR spectra of TGDDM-PDA epoxy formed in the absence of any polyimide.

Figure 7. FTIR-ATR spectra of TGDDM-PDA epoxy formed in the presence of polyimide.

Figure 8. Fraction of insoluble gel in polyimide/epoxy (PI/TGDDM-PDA) and polyimide/diamine (PI/PDA) films, as a function of the number of diamine units available for cross-linking reactions.

this increase in free volume arises only from a reduction in the number of cross-links forming. FDA has the highest molecular weight, so it provides fewer diamine moieties at any given wt % epoxy for cross-linking. There is no evidence that the bulkiness of the FDA molecule results in a more open structure as we had hoped. Conversely, polyimide films containing only PDA diamine show considerably greater free volume, indicating that this approach to cross-linking retains more of the original polyimide structure. Microstructural Analysis. In Figure 10, data obtained from wide-angle X-ray diffraction analysis of 6FDA-TMPDA/ TGDDM-FDA films is presented. It is evident that as the amount

of epoxy is increased, the WAXD peak intensity reduces; however, no change in the position of the peak is observed. This suggests that while the average spacing between the polymer chains is unchanged, the number of chains in an ordered position is reduced. The WAXD spectra for films prepared with the other two different diamine structures are comparable to Figure 10, indicating similar changes in the microstructure. These results are consistent with previous results obtained with diamine cross-linking of polyimides by both our own group14 and others.7 The density, FFV, and WAXD results together indicate that as cross-linking occurs, the proportion of the polymer that exists in an ordered and open structure, characteristic of a glassy polyimide, reduces and the proportion of denser, more amorphous structure increases. Thermal Analysis. The temperature at which a given compound degrades depends on the thermal stability of its constituent functional groups. Often, this characteristic is quantified in terms of the degradation temperature, TD: the temperature at which the rate of mass loss is highest. The TGA thermograms for these systems are provided in Figure 11 and degradation temperatures are given in Table 2. While the pure polyimide film is stable to around 480 °C, the pure epoxy sample shows a TD of 361 °C. As expected, the 0.133 mol/mol of PI/epoxy sample (0.133 mol diamine/mol repeat unit) lies between these two extremes, consistent with the formation of epoxy structures. However, at higher temperatures, a second significant increase in mass loss also occurs in this sample, suggesting that an amount of polyimide in the

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Figure 9. Fractional free volume of PI/epoxy films plotted against (a) mass concentration of the epoxide/diamine precursor and (b) the number of diamine units available for cross-linking reactions.

Figure 10. Wide-angle X-ray diffraction analysis of polyimide/epoxy (6FDA-TMPDA/TGDDM-FDA) samples. Figure 12. Logarithmic plot of the permeability of various gases versus reciprocal FFV. Measurements were conducted at 1000 kPa and 35 °C and permeability is calculated on a fugacity basis. Data for all polyimide/epoxy films are presented as solid symbols. Experimental data for carbon dioxide permeability in polyimide/diamine films and literature results obtained by Lin et al.8 are also shown as open symbols. Error bars detail the 95% confidence interval for each point, based on a pooled standard deviation determined from replicate measurements. Table 2. Degradation Temperatures (TD) for Pure Epoxy, Pure Polyimide (PI), Polyimide/Epoxy, and Polyimide/Diamine Variants as Measured by Thermogravimetric Analysis

sample

Figure 11. Thermogram of polyimide/epoxy film variants.

sample may not have cross-linked with the epoxy material, and so is stable to a higher temperature. Thermal analysis of polyimide/diamine samples containing 0.147 molPDA molPI_Repeat_Unit-1 revealed a degradation temperature even greater than the polyimide control (result not shown). This gives further weight to the assertion that inclusion of TGDDM epoxide in a polyimide/diamine system can give rise to different cross-linking mechanisms than when diamine is added alone. The polyimide/epoxy films were also characterized using differential scanning calorimetry. From the pure polyimide scan, a Tg of 422 ( 1.5 °C was determined, which compared favorably to the literature values, range 420-424 °C.26,27 The scans for the polyimide/epoxy films were more complex and it is difficult to clearly define a glass transition temperature. This reflects

TGDDM-TMPDA epoxy PI/TGDDM-TMPDA epoxy PI/TGDDM-PDA epoxy PI/TGDDM-FDA epoxy PI PI/PDA

cross-linker concentration (mol/molPI_Repeat_Unit) 0 0.101 0.133 0.147 0.147

TD (°C) 361 416 419 433 >460 >480

both the low-temperature degradation reactions of the epoxy compounds and the complexity of the grafted polyimide structure present in these systems. Such cross-linked networks have previously been shown to exhibit broad and complex glass transitions.28 Gas Transport. Figure 12 shows the permeability of several gases for a range of polyimide/epoxy structures at 1000 kPa upstream pressure. All three diamine structures are included within the above data. The range of CO2 permeabilities for 6FDA-TMPDA at 1000 kPa upstream pressure and 35 °C reported in the literature is 440-612 Barrer6,8,18 on a pressure basis (∼463-644 Barrer on a fugacity basis). In the present work, we obtain a value of 703 Barrer on a fugacity basis. This

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Figure 13. Permeability versus selectivity for the epoxy systems examined in this paper. Results generally move in a direction parallel to Robeson’s upper bound.33 Data for polyimide/PDA films is sourced from Powell et al.14

Figure 14. CO2 permeation through polyimide/TGDDM-TMPDA at 35 °C: (a) Permeability data calculated on a fugacity basis and (b) the same permeability data normalized by the value recorded at a feed pressure of 1000 kPa.

high value can be attributed to the use of moderate membrane curing conditions leading to a high fractional free volume. It is evident from Figure 12 that our data is consistent with the data of Lin et al.8 when this fractional free volume is considered. As the epoxy concentration is increased, the fractional free volume falls. This results in an exponential fall in permeability, as given by free volume theory, eq 1,29

P ) Ae(-B/FFV)

(1)

where P is the permeability, FFV is the fractional free volume, and the parameters A and B are related to the penetrant type. Data for polyimide/diamine films also fit this relationship. As these films undergo much smaller changes in fractional free volume, there is also a smaller decrease in permeability relative to the pure polyimide film. As expected from free volume theory, the loss of permeability upon epoxy addition is accompanied by an overall increase in selectivity. Indeed, while there is some scatter in the data, Figure 13 shows that increasing the epoxy concentration leads to polymer performance generally moving in a parallel direction to Robeson’s upper bound33 consistent with this theory.34 The permeability of pure CO2 through PI/TGDDM-TMPDA epoxy systems at 35 °C and in the pressure range 330-3000 kPa is presented in Figure 14a. The loss of permeability as epoxy is added is clearly evident. However, from the curve shape it appears that the epoxy addition has suppressed plasticization, as the permeability for films containing 0.133 mol/mol no longer increases at high fugacities of this condensable gas.

The permeability data for 12 wt % epoxy (0.101-0.147 mol/ mol) is plotted relative to the value at 1000 kPa in Figure 14b. Also plotted on the same scale is the data for the pure polyimide film and for a polyimide/diamine film. At this finer scale, it is clear that even when epoxy is present, the permeability does indeed fall a little over the initial pressure range as the solubility of CO2 in the polymer increases. For the polyimide film, the permeability increases again at high pressures due to plasticization effects. However, for both the polyimide/diamine and polyimide/epoxy membranes, there is no turning point in the curve, confirming that the membrane swelling effects associated with plasticization have been significantly depressed. This is consistent with the work of Wind et al.35 who show that crosslinking decreases polymer swelling and thus stabilizes the CO2 diffusion coefficient under high CO2 pressures. Further to this data, it is useful to compare the plasticization resistance of the film over extended time periods at high CO2 pressures. It is evident from Figure 15 that the inclusion of epoxy material into the membrane has resulted in smaller increases in CO2 permeability during a 48 h CO2 exposure period. This implies that incorporation of epoxy material into the membrane has increased the film’s plasticization resistance. The effect on polymer relaxation rate was also investigated. Since relaxation is thought to be a first-order process,35 the relaxation rate constant can be determined from the gradient of a semilogarithmic plot of the normalized permeability against exposure time to high-pressure carbon dioxide. As shown in Figure 16, there is no evidence that the epoxy addition has

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Research Centre of the Australian Research Council, and this support is also greatly appreciated. Literature Cited

Figure 15. Plasticization of polyimide-based films by CO2 at 3100 kPa and 35 °C. The permeability at time t (Pt) is normalized by the permeability at time t ) 1 h(P1).

Figure 16. Relaxation of PI/epoxy systems at 35 °C when exposed to CO2 at 3100 kPa. P1 represents the initial permeability measured after 1 h. P48 represents the final permeability measured after 48 h. A semilogarithmic plot of (1 - (Pt - P1)/(P48 - P1) versus time should be linear if the change in permeability is rate-controlled by a first-order process such as polymer relaxation.35

altered the relaxation rate. While somewhat surprising, this result is consistent with our earlier results with diamine systems.14 Conclusions A range of epoxide/diamine systems have been added to 6FDA-TMPDA polyimide, creating graft network membranes. The density and microstructure of these membranes have been examined and these indicate significant epoxy formation within the network. This was confirmed using infrared spectroscopy. Epoxy incorporation leads to the partial collapse of the open structure of the glassy polyimide, leading to a more dense, less ordered structure. These effects were found to be independent of diamine structure. As the epoxy concentration was increased, gas permeability decreased, while selectivities improved, consistent with established free volume theory. Additionally, the cross-linked structure resulted in a suppression of carbon dioxide-induced plasticization. Such plasticization resistance, however, is similar to that exhibited by polyimide films containing diamine only. Furthermore, such diamine cross-linked films were observed to have significantly higher permeabilities, with similar selectivities. Hence, they provide more attractive gas separation properties than the polyimide/epoxy films studied in the present work. Acknowledgment Financial support for this project has been provided by the CRC for Greenhouse Gas Technologies (CO2CRC) and this support is gratefully acknowledged. Infrastructure support is also provided by the Particulate Fluids Processing Centre, a Special

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ReceiVed for reView May 15, 2007 ReVised manuscript receiVed August 16, 2007 Accepted September 12, 2007 IE070689H