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Highly CO2-selective gas separation membranes based on segmented copolymers of poly(ethylene oxide) reinforced with pentiptycene-containing polyimide hard segments Shuangjiang Luo, Kevin A Stevens, Jaesung Park, Joshua D Moon, Qiang Liu, Benny D. Freeman, and Ruilan Guo ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.5b11355 • Publication Date (Web): 06 Jan 2016 Downloaded from http://pubs.acs.org on January 12, 2016

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ACS Applied Materials & Interfaces

Table of Contents Graphic

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Highly CO2-selective gas separation membranes based on segmented copolymers of poly(ethylene oxide) reinforced with pentiptycene-containing polyimide hard segments Shuangjiang Luo,† Kevin A. Stevens,‡ Jae Sung Park,‡ Joshua D. Moon,‡ Qiang Liu,‡ Benny D. Freeman,‡ Ruilan Guo†* †

Department of Chemical and Biomolecular Engineering, University of Notre Dame, Notre Dame, IN 46556, United States ‡ Center for Energy and Environmental Resources, Department of Chemical Engineering, Texas Materials Institute, The University of Texas at Austin, Austin, TX 78758, United States * Corresponding author: [email protected], +1-574-631-3453 (tel), +1-574-631-8366 (fax) KEYWORDS: CO2-selective gas separation membrane, poly(ethylene oxide), pentiptycene, segmented copolymer ABSTRACT: Poly(ethylene oxide) (PEO)-containing polymer membranes are attractive for CO2-related gas separations due to their high selectivity towards CO2. However, the development of PEO-rich membranes is frequently challenged by weak mechanical properties and a high crystallization tendency of PEO that hinders gas transport. Here we report a new series of highly CO2-selective, amorphous

PEO-containing

segmented

copolymers

prepared

from

commercial

Jeffamine®

polyetheramines and pentiptycene-based polyimide. The copolymers are much more mechanically robust than the non-pentiptycene containing counterparts due to the molecular reinforcement mechanism of supramolecular chain threading and interlocking interactions induced by the pentiptycene structures, which also effectively suppresses PEO crystallization leading to a completely amorphous structure even at 60% PEO weight content. Membrane transport properties are sensitively affected by both PEO weight content and PEO chain length. A nonlinear correlation between CO2 permeability with PEO weight content was observed due to the competition between solubility and diffusivity

contributions,

whereby

the

copolymers

change

from

being

size-selective

to

solubility-selective when PEO content reaches 40%. CO2 selectivities over H2 and N2 increase 1 ACS Paragon Plus Environment

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monotonically with both PEO content and chain length, indicating strong CO2-philicity of the copolymers. The copolymer film with the longest PEO sequence (PEO2000) and highest PEO weight content (60%) showed a measured CO2 pure gas permeability of 39 Barrer, and ideal CO2/H2 and CO2/N2 selectivities of 4.1 and 46, respectively, at 35 oC and 3 atm, making them attractive for hydrogen purification and carbon capture.

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1. Introduction Compared with conventional gas separation techniques such as pressure swing adsorption and cryogenic distillation, membrane-based gas separation has attracted substantial attention during the past decades due to its inherent advantages such as high energy efficiency, small footprint, relatively low cost, ease of operation and simple maintenance.1-3 Since the late 1970s, polymer membranes have found wide spread applications for commercial gas separation processes including nitrogen enrichment from air (O2/N2), hydrogen purification and recovery from ammonia purge and refinery processes (e.g., H2/N2 and H2/CH4), syngas ratio adjustment (H2/CO), natural gas sweetening (CO2/CH4), and dehydration (H2O/air).1, 3-4 Glassy aromatic polymers such as polyimides are attractive gas separation membrane materials due to their strong size-sieving ability as well as excellent mechanical properties.1 However, they are not optimal for performing some CO2-related separations such as hydrogen purification (CO2/H2), when selective transport of large gas molecules (i.e., CO2) over small gas molecules (i.e., H2) can be desirable to reduce or eliminate the costly processes of H2 recompression by keeping the smaller gas in the high-pressure retentate.5 In these cases, membranes featuring solubility controlled transport are needed. Compared to the diffusivity controlled transport process, the solubility-selective process is substantially thermodynamic, and the affinitive polymer-penetrant interactions as well as the condensability of the penetrants dominate the transport process.6 For polar or acid gases (e.g., CO2), polymers containing polar moieties, such as poly(ethylene oxide) (PEO) with ether groups, are generally more selective to CO2 than to other gas molecules, presumably due to favorable dipole-quadrupole interactions between CO2 and the polar ethers.7-10 However, due to the poor film-forming ability of pure PEO, efforts to fabricate CO2-selective PEO-containing membranes have been largely focused on copolymerizing PEO with glassy polymers,11-17 blending PEO with a film-forming polymer,18-20 or use of UV-cross-linked poly(ethylene glycol)-based acrylates and related materials.21 High PEO content and long PEO sequences are usually preferred to achieve high CO2 permeability and selectivity because well-connected PEO domains provide pathways for fast and selective CO2 transport. However, the high crystallinity of long PEO chains counteracts the fast 3 ACS Paragon Plus Environment

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permeation rates,7 and the issues of mechanical and thermal instability become more substantial when PEO content is high. Covalently crosslinked PEOs represent a promising strategy to circumvent the aforementioned dilemma in that crosslinking effectively disrupts crystallization, which allows high permeability and high selectivity.21-24 However, these crosslinked PEOs are highly hydrophilic, which can contribute to challenges in forming ultra-thin or asymmetric films from these polymers. Moreover, pure PEO membranes, even highly crosslinked, have limited thermal and mechanical stability.25 Therefore, highly amorphous PEO-rich membranes that are mechanically strong and thermally stable are desirable for CO2 separations. In our recent studies of iptycene-containing polyimide gas separation membranes,26-27 we found that the molecular cavities of the iptycene units are accessible to the neighboring side groups and polymer chains, which provides a useful supramolecular mechanism to finely tune the chain packing and organization in polymers. This finding provides a strong incentive to incorporate iptycene structure into PEO-rich copolymers to provide molecular-level reinforcement and suppress PEO crystallization via strong supramolecular interactions between the iptycene units and PEO sequences. As depicted in Figure 1, flexible PEO sequences may thread through the clefts of iptycene units (pentiptycene in this case) delineated by the arene blades and get confined or interlocked via physical entanglements. As demonstrated in several triptycene-based polycarbonates and polyesters, significantly improved mechanical properties have been achieved via this supramolecular chain threading and interlocking mechanism.28-30 For example, a triptycene-containing polyester exhibited a nearly three-fold increase in Young’s modulus and strength, and a more than twenty-fold increase in strain to failure when compared with the mechanical properties of a non-triptycene polyester.28 As such, mechanical interlocking of iptycene units allows facile molecular level manipulations to achieve tailored membrane morphologies and significantly enhanced mechanical properties. Additionally, chain threading and interlocking with iptycene units may effectively interrupt runs of PEO sequences, thereby frustrating PEO crystallization and facilitating faster diffusion.



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Figure 1. Schematic of a PEO chain (blue dotted tube) threading through one of the molecular clefts of pentiptycene unit. In this study, we have developed robust CO2-selective segmented copolymer membranes from commercial Jeffamine® polyetheramines and pentiptycene-based polyimide as the hard segments. Pentiptycene features a rigid, H-shaped structure constructed by five benzene rings fused together, which has more and larger open molecular cavities than triptycene structure available for PEO chain threading.26-27 PEO molecular weight and PEO weight content were systematically adjusted in these copolymer membranes, which permitted tailoring of membrane physical properties and morphology, fractional free volume, and consequently gas transport properties, as well as elucidation of fundamental structure-property relationships for these new PEO-rich segmented copolymers.

2. Experimental Section Materials. 6,13-Bis(4-amino-2-trifluoromethylphenoxy)pentiptycene (PPDA) was synthesized according to our recently published procedure.27 Polyetheramines with molecule weights of 600, 900, 2000 g⋅mol−1 (i.e., Jeffamine® ED-600, ED-900 and ED-2003) were obtained from Huntsman Co. and used as received. These polyetheramines are referred as PEO600, PEO900 and PEO2000 in this paper for simplicity. 4,4’-Hexafluoroisopropylidene bisphthalic dianhydride (6FDA, ˃99.0%) was purchased from Akron Polymer Systems and dried at 160 oC in a vacuum oven for 16 hours prior to use. Anhydrous N,N-dimethylacetamide (DMAc, 99.8%), anhydrous pyridine (99.8%) and acetic anhydride (˃98.0%) were purchased from Sigma-Aldrich (USA) and used without further purification. Synthesis of segmented PEO-containing copolymers. Segmented copolymers with precisely controlled PEO weight content were synthesized via condensation polymerization between PPDA, 5 ACS Paragon Plus Environment

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polyetheramine and 6FDA using chemical imidization method (Scheme 1). Synthesis of a copolymer with 40 wt% of PEO2000 is provided as an example: a mixture of PPDA (0.8614 g, 1.10 mmol) and PEO2000 (1.0124 g, 0.51 mmol) was dissolved in anhydrous DMAc (11 mL) in a flame-dried, three-necked flask under nitrogen atmosphere at room temperature. After dissolution, the flask was cooled to 0 oC in an ice water bath, then a stoichiometric amount of 6FDA (0.7150 g, 1.61 mmol) was added into the flask with stirring. The mixture was warmed gradually to room temperature and stirred overnight to obtain a viscous PEO-poly(amic acid) solution. To chemically convert it to imide structure, acetic anhydride (0.9 mL) and pyridine (0.9 mL) were added into the flask and stirred for another 24 h. The final solution was precipitated in an excess amount of methanol to give fibrous copolymer products, which were collected by filtration, washed with methanol and then dried at 150 oC under vacuum overnight. Synthesis of other copolymers followed the same procedure except that different ratios of PPDA/PEO-diamine or PEO-diamine with different chain length were used to systematically vary the PEO content and sequence length in the resulting copolymers. The copolymers are named as pent-PI-PEOx (y%), where x represents the PEO sequence length and y is the PEO weight content. For example, pent-PI-PEO900 (40%) copolymer has 40% of PEO by weight, which has a chain length of 900 g⋅mol−1.

Scheme 1. Synthesis of segmented pent-PI-PEO copolymers via condensation copolymerization of PPDA, Jeffamine® polyetheramine and 6FDA (Note: the structure of Jeffamine® polyetheramines is simplified for clear view). Fabrication of dense films. Dense thin films of the copolymers were prepared by solution casting 6 ACS Paragon Plus Environment


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method. Predetermined amount of polymer was dissolved in DMAc (~ 7% w/v, g⋅mL−1) and filtered through a 0.45 µm PTFE syringe filter before casting on a clean, flat, leveled glass plate. The solution was heated under an infra-red lamp at about 60 oC overnight to allow slow evaporation of the solvent and form a film. The obtained isotropic film was further heated at 160 oC overnight under vacuum to completely remove the casting solvent. All the films were solvent free as confirmed by thermo-gravimetric analysis (TGA). Film samples for permeation measurements were epoxied (Devcon, No 145250, Danvers, MA, USA) to a brass shim stock disk with a hole drilled in it to control the exposed area. Films in the glassy state at 35 °C were supported by a filter paper backing to prevent contamination, while rubbery films were supported by a Whatman Anodisc (No. 6809-6022) to provide a flat surface and avoid deformation against the filter paper. Effective area for gas permeation measurements was determined using a scanner (LiDE120, Canon) and ImageJ software, and the film thickness (typically 60-120 µm) was measured using a digital micrometer. Characterization methods. Fourier transform infrared spectra of the polymer films were recorded on a Jasco FT/IR-6300 spectrometer in attenuated total reflection mode (ATR-FTIR) with a resolution of 4 cm‒1 and 64 scans. 1H NMR spectra were recorded on a Bruker 500 spectrometer using deuterated dimethyl sulfoxide (DMSO-d6) as a solvent. Molecular weight and molecular weight distribution were measured with size exclusion chromatography (SEC, Waters GPC System) at 35 oC using polystyrene standards and DMF as the eluent. The DMF SEC was equipped with Polymer Standards Service (PSS) columns (guard, 105, 103, and 102 SDV columns) with DMF flow rate of 1 ml·min-1, and a differential refractive index (RI) detector (Wyatt Technology, Optilab T-rEX) using PSS WinGPC 7.5 software. Polymer thermal stability was evaluated by thermo-gravimetric analysis (TGA), using a TGA Q500 instrument (TA Instruments) at a heating rate of 10 oC⋅min−1 with a nitrogen purge of 60 mL⋅min−1. Differential scanning calorimetry (DSC) measurements were conducted on a DSC Q2000 (TA Instruments) with liquid nitrogen cooling system at a heating rate of 10 oC⋅min−1 and cooling rate of 20 C⋅min−1 with a nitrogen purge of 50 mL⋅min−1. Samples were evaluated in the range of -100 to 300 oC,

o


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and the second heating profile was reported. Glass transition temperature (Tg) was determined using the automatic mode of TA Universal Analysis software. In both TGA and DSC measurements, samples were in thin film state. The water contact angles of the films were determined by a goniometer (Rame-Hart, model 250) to assess their surface hydrophilicity. Sessile contact angle titrations were conducted more than five times for each film with 4 µL deionized (DI) water drops in air, and the averaged values were reported. Membrane surface morphology was captured by atomic force microscopy (AFM) on a Park System XE-70 instrument in tapping mode at room temperature with a silicon AFM probing tip (TESPA, Bruker Nano). At least five images were taken at different positions for each sample, and both the height image and phase image were obtained using WSxM 5.0 software.31 Mechanical properties of the films were tested using a Shimadzu Autograph AGS-X 500N tensile testing instrument at room temperature. At least three films were tested for each sample and the average values were reported. All films were prepared with an initial width of 5 mm, gauge length of 22 mm and thickness of 80-120 µm. The uniaxial tensile speed was 1 mm·min-1 (ASTM D-1708-13). Tensile strength was reported as the maximum stress observed during the uniaxial tensile tests, and the elongation was read at the point where the sample failed. Wide-angle X-ray diffraction (WAXD) data were recorded using a Bruker D8 ADVANCE DAVINCI diffractometer with Cu Kα radiation (wavelength λ = 1.5418 Å) operated at 40 mA and 40 kV. The scan speed and step size were 5 seconds per step and 0.02o per step, respectively. The densities of PEO-containing copolymer films were measured using an analytical balance (ML204, Mettler Toledo) equipped with a density kit using the buoyancy method at room temperature. Hexanes was chosen as the medium because these films showed negligible absorption (< 1 wt %) of hexanes after being soaked in hexanes for 24 h at room temperature. The density of hexanes was determined using a pycnometer at room temperature. Fractional free volume (FFV) of the copolymers was estimated using the following equation:

FFV =

V0 −1.3Vw V0

(1)



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where V0 is the specific volume of the films obtained from the density measurement, and Vw is the van der Waals volume determined using Bondi’s group contribution method.32-33 In the case of copolymers, Vw was estimated by: Vw = x1·Vw1 + x2·Vw2, where x1 and x2 are the molar fractions of the PEO segments and polyimide segments, respectively; Vw1 and Vw2 are the respective van der Waals volumes for the two components. Pure gas permeabilities of H2, CH4, N2, O2 and CO2 were measured in that order at 35 oC using a custom-built permeation system based on the constant-volume, variable-pressure method.34 All films were masked as described previously and degassed overnight before measurements. Upstream pressure was maintained at predetermined pressures (i.e., 3.0, 6.4, 9.8, 13.0 and 16.6 atm), while the downstream pressure increase was recorded over time. Pure gas permeabilities were determined as follows:34

P = 1010

Vd l  dp dp  ( )ss − ( )leak   pupTRA  dt dt 

(2)

where P (Barrer, 1 Barrer = 10‒10 cm3(STP)·cm/cm2·s·cmHg) is the gas permeability, Vd is the downstream volume (cm3), l is the film thickness (cm), pup is the upstream pressure (cmHg), A is the effective membrane area (cm2), (dp/dt)ss is the steady-state pressure increment in the downstream (cmHg/s), (dp/dt)leak is the leak rate of the system (cmHg/s), T is the test temperature (K), and R is the gas constant (0.278 cm3·cmHg/cm3(STP)·K). The ideal selectivity (αA/B) for two different gases A and B was defined as the ratio of pure gas permeability of these two gases:

α A/ B ≡

PA PB

(3)

3. Results and Discussion Copolymer Synthesis and Characterization. Copolymers with systematically varied PEO weight content and chain length were prepared via two-step condensation polymerization (Scheme 1).

A

diamine mixture consisting of pentiptycene-based diamine (PPDA) and Jeffamine® polyetheramine (PEO600, PEO900 or PEO2000) in a predetermined molar ratio reacted with 6FDA dianhydride (in equal molar amount of diamines) to obtain the poly(amic acid) intermediate, which was 9 ACS Paragon Plus Environment

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cyclodehydrated chemically using acetic anhydride and pyridine to produce the final imide copolymers. By adjusting the molar ratio of PPDA to PEO-diamine, the weight content of PEO in the copolymers was systematically varied in the range of 20–60%. For given PEO weight fraction, copolymers with varying PEO sequence length were prepared by using PEO600, PEO900 or PEO2000.

Figure 2. 1H NMR spectra (DMSO-d6) of pent-PI-PEO900 copolymers with various PEO content. The dashed line framed region (3.0 – 3.6 ppm) is enlarged to show the intensity change of peak j. The completely imidized structures of the copolymers were confirmed using 1H NMR and ATR-FTIR. Figure 2 shows the 1H NMR spectra of pent-PI-PEO900 copolymer series with various PEO weight content, wherein the peak assignment confirms the anticipated chemical structures of the copolymers. 1H NMR spectra of other copolymers are similar and are not shown here. The intensity of PEO peaks (i.e., peaks k, l and j) increases with increasing PEO weight percent, suggesting the copolymerization was successful. Moreover, PEO content can be quantified by comparing the peak intensities of the PEO protons (e.g., peak j) with the pentiptycene protons (e.g., peak c). As shown in Table 1, the PEO contents estimated from 1H NMR spectra are very close to the target values, indicating successful control of the copolymer composition and high reactivity of all monomers 10 ACS Paragon Plus Environment


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towards copolymerization. ATR-FTIR spectra (normalized by referencing imide carbonyl band at 1726 cm-1) of all the copolymers (Figure 3) show the characteristic asymmetric imide carbonyl stretching at 1784 cm-1 and symmetric imide carbonyl stretching at 1726 cm-1. No carbonyl stretching of the poly(amic acid) peak was observed in the range of 1620-1680 cm-1, confirming complete imidization. Additionally, the intensity of the PEO peaks (e.g., 2870 cm-1) increases with increasing PEO weight content, as expected.

Figure 3. ATR-FTIR spectra of PEO-containing copolymer membranes. Molecular weight and molecular weight distribution of the copolymers were obtained from SEC measurements. As shown in Table 1, all copolymers show high molecular weight with polydispersity in the range of 2.0-3.7, indicating good reactivity of Jeffamine® polyetheramines and PPDA towards polycondensation reaction. High molecular weight also aids in preparing defect-free thin films via conventional solution casting method. Table 1. Summary of PEO weight content, molecular weight and thermal properties of pent-PI-PEO copolymers. Copolymer

PEO

Mn

PDI

Tg

Tg, Fox


Tg, G-T

Td, 5%

Td2

Char

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wt% a

(g⋅mol−1)

(Mw/Mn)

(oC)

(oC)b

(oC)c

(oC)d

(oC)e

Yield (%)f

pent-PI-PEO600 (20%)

18.8

94, 900

2.94

188

—

—

388

530

58


18.8

58, 600

2.04

202

—

—

383

530

61


19.5

73, 500

2.02

204

186

99

382

530

63


38.0

117, 800

3.71

56

—

—

343

530

31


40.9

119, 800

2.45

58

—

—

364

520

37


37.8

46, 300

2.22

65g

88

19

350

520

45


46.8

188, 800

3.58

-18

54

-3

369

520

37


58.8

80, 800

2.41

-20

—

—

364

520

13


56.8

53, 800

2.16

-35

22

-21

350

520

27

a

Actual PEO weight content determined from 1H NMR. b Glass transition temperature of PEO2000 series estimated by the Fox equation using the actual PEO weight contents. c Glass transition temperature of PEO2000 series estimated by Gordon-Taylor equation. d Td, 5%: decomposition temperature at 5% weight loss. e Td2: onset decomposition temperature in the second stage. f Residual weight retention after heated to 600 oC in nitrogen atmosphere. g This value is an estimation as the glass transition is very subtle in DSC curve.

DSC measurements were conducted to investigate the crystallization tendency of PEO component and the glass transition temperature of the copolymers (Table 1 and Figure 4a). For each of the copolymers, only one Tg was detected, suggesting the two components formed a compatible or random copolymer system with no appreciable phase separation. Incorporation of soft PEO sequences into the copolymer backbone greatly reduced the glass transition temperature relative to that of 6FDA-PPDA polyimide (374 oC).27 For the pent-PI-PEO2000 series, the Tg dropped by 170 oC with only 20% PEO incorporation, and further decreased from 204 oC to -35 oC as PEO weight content increased from 20% to 60%, suggesting significantly increased chain flexibility. Similar trends were reported in several other PEO-containing copolymer systems.12,

14-15, 35-37

It is important to note that the copolymers

containing 40% PEO are still glassy at 35 oC (the temperature used in permeation tests) regardless of PEO chain length, and the Tg is much higher than the non-iptycene containing PEO-based copolymers reported in the literature having similar PEO contents.12, 14, 35-36 This observation is consistent with supramolecular chain threading associated with the pentiptycene structures (Figure 1), which induces 12 ACS Paragon Plus Environment


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strong molecular confinement so as to greatly reduce PEO chain mobility and results in glassy copolymers even with 40% PEO content. The Tg values of the pent-PI-PEO2000 series were compared with theoretical values predicted by both Fox equation (4)38 and Gordon-Taylor equation (5)39:

1 w1 w2 = + Tg Tg,1 Tg,2 Tg =

(4)

wT 1 g,1 +kw2Tg,2

(5)

w1 + kw2

where w1 and w2 are the weight fractions of PEO and polyimide, respectively; Tg,1 (PEO2000) and Tg,2 (6FDA-PPDA polyimide) are 209 K and 647 K, respectively; k is a constant parameter which represents the unequal contributions of components to the blend, and its value (0.143) was determined from k ≈ (V2Tg,1 / V1Tg,2).40 The trend of the measured Tg values match well with that predicted by the Fox equation and Gordon-Taylor equation, specifically, Fox equation predicted well at low PEO contents (i.e., 20% and 40%) while Gordon-Taylor equation has a better prediction at high PEO content (i.e., 50% and 60%). As is known, both Fox equation and Gordon-Taylor equation predict a rather simplified thermodynamic relation, which gives good predictions on compatible blends and statistical copolymers only when the Tgs of the two components are not too different, or the product of the heat capacity and Tg are very close for the constituents.41 In this study, soft PEO and hard 6FDA-PPDA polyimide sequences have significantly different thermal characteristics with over 430 oC difference in their Tgs. As such, neither equation was able to provide good predictions for the copolymers in this study, especially when non-conventional supramolecular inter-chain interactions play a significant role in influencing local chain dynamics. For all copolymers, no endothermic/exothermic peaks were observed in the range of 20 oC to 50 oC, which is the characteristic melting/crystallization temperature range of PEO crystals12. The suppression of PEO crystallization is due to the combined effects of chain packing disruption by the hard polyimide segments and the above-mentioned chain threading interactions, both of which hinder the free movement of PEO segments, break the runs of PEO sequences, and prevent the formation of PEO crystals.28 13 ACS Paragon Plus Environment

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Figure 4. (a) DSC curves of pentiptycene-PEO-containing copolymers as functions of PEO weight content and chain length (names of the copolymers next to each curve are simplified by only showing the PEO chain length and weight content); (b) TGA thermograms of pent-PI-PEO2000 copolymer series with different PEO weight content. The effect of PEO chain length on the glass transition temperature of the copolymers, however, is more complicated, which shows different trends for different PEO content. At low to medium PEO content (i.e., 20% and 40%), copolymers with longer PEO sequences had higher Tg values; at high PEO content (i.e., 60%), the opposite trend was observed. These observations may be rationalized by the combined effects of limited overall chain mobility induced by hard segments and the chain threading effect of pentiptycene structures. Specifically, at the same PEO weight content, the longer the PEO sequence, the higher the molar ratio of the hard polyimide segments, which contributes more towards restriction of PEO segment mobility. At low to medium PEO content, the PEO segments are relatively separated or isolated by the hard pentiptycene-polyimide segments. The chain threading effect may be more significant for long PEO sequences than for short PEO ones, resulting in higher Tg values of the copolymers containing long PEO sequences. However, when the soft PEO segments become dominant (60%) in the copolymers, long PEO chains, even partially threading through pentiptycene units, are less restricted, leading to lower Tgs than the short-PEO-sequence copolymers at a given PEO content. Thermal stability of PEO-containing copolymer films was evaluated by TGA, and the results are shown in Table 1 and Figure 4b. All the TGA thermograms show a two-stage degradation profile. The 14 ACS Paragon Plus Environment


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first stage starts around 340 oC, and the weight loss values during this stage match well with the PEO weight content in the copolymers, consistent with the anticipated composition of these copolymers. The second stage is due to the degradation of the aromatic pentiptycene-polyimide component at around 530 oC, and the char yields at 600 oC under N2 show an inverse relationship with PEO content as expected.36, 42 Incorporating pentiptycene-based imide structures greatly improved the thermal stability of the copolymers. It is also noted that there is no premature weight loss due to moisture or the residual casting solvent, indicating that all films were solvent free after the high temperature vacuum drying process. Membrane Morphology and Properties. For multicomponent membranes, the detailed microdomain morphology, such as the shape and interspatial arrangement (i.e., the connectivity) of the respective phases, is critical in determining the membrane properties, which are sensitively influenced by the copolymer composition, the length of segmental sequences, and the interactions between the components.12, 14, 43-44 Surface morphology of the copolymer films was imaged with AFM. AFM phase images along with the height images of the pent-PI-PEO2000 series of films with different PEO content are shown in Figure 5. The dark regions in the AFM phase images (top row in Figure 5) represent the soft PEO components, and the bright areas indicate the hard pentiptycene-polyimide components. As shown, no distinct phase-separated pattern can be defined from the phase images for the copolymer films, which supports the hypothesis that the hard and soft segments are randomly/statistically distributed throughout the film. This observation is consistent with the DSC results, where only one glass transition temperature was detected, and no melting/crystallization peaks were observed even with 60% PEO content. Increasing the PEO content neither led to large distinct domains nor induced phase separation, unlike what is typically observed for PEO-containing copolymers without pentiptycene structures.12,

43-44

This is likely because penetration of soft PEO sequences into the hard

pentiptycene-polyimide segments via chain threading might blur the phase boundary. Additionally, the chain length of the PEO sequences in these copolymers is relatively short (i.e., the highest MW was 2000 g⋅mol−1), which does not favor sharp phase separation. It was reported in several studies that high 15 ACS Paragon Plus Environment

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molecular weight PEO (i.e., more than 10,000 g⋅mol−1) was necessary to decrease the inter-penetration between soft and hard segments leading to distinct phase separation.12, 43-44 AFM height images (bottom row in Figure 5) show that all films have smooth surfaces with roughness values of less than 10 nm (film thickness of ~ 60 μm).

Figure 5. AFM phase images (the top row) and height images (the bottom row) of pent-PI-PEO2000 membranes with PEO content of a) 20%, b) 40% and c) 60%. Surface hydrophilicity of the copolymer membranes was evaluated via contact angle measurements as a function of PEO content and PEO chain length. As shown in Table 2 and Figure S1, contact angle decreased almost linearly by about 20o when the amount of hydrophilic PEO component increased from 20% to 60%, indicating gradually increased surface hydrophilicity. At the same PEO content, copolymer films prepared from different PEO lengths have very similar contact angle. This observation suggests that film surface hydrophilicity is mainly determined by the PEO content. Previous studies demonstrated that supramolecular chain threading and interlocking induced by triptycene structures could significantly enhance the mechanical properties of corresponding iptycene-containing copolymers.28-30 Mechanical properties of the pent-PI-PEO copolymer films in terms of Young’s modulus, elongation at failure and tensile strength are listed in Table 2. As a general observation, pent-PI-PEO copolymers are much more mechanically robust than previously reported 16 ACS Paragon Plus Environment


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PEO-containing copolymers having similar PEO content.12, 35 For example, the Young’s modulus of the pent-PI-PEO2000 (60%) copolymer is more than 15 times higher than that of 6FDA-ODA-PEO1 (60) reported in the literature12, which has a very similar chain structure with the same PEO content and PEO sequence but contains no pentiptycene units. These results that threading of flexible PEO sequences through the clefts of the pentiptycene units and supramolecular interlocking between pentiptycene units effectively provide strong molecular-level reinforcement leading to significantly improved Young’s moduli.28-30 Within the pent-PI-PEO copolymer series in this work, both Young’s moduli and tensile strength decreased by an order of magnitude as PEO content increased from 20% to 60%, and the copolymers became thermoplastic elastomers at high PEO contents with very large deformations to break indicating the mechanical strength of the copolymers is mainly determined by the PEO content. The effect of PEO sequence length on the mechanical properties, however, did not follow a clear trend, although the Young’s moduli seemed to decease slightly with PEO chain length for a given PEO content. In PEO-containing copolymer membranes, the role of polar PEO sequences is mainly to selectively interact with polar gases such as CO2 in the feed mixture to enhance the solubility selectivity towards the polar gases. Ideally, based on the solution-diffusion model for dense membranes,45 the enhanced solubility contribution from the PEO segments should lead to improved CO2 permeability if the diffusivity component is not adversely affected by the PEO segments. In the case of pentiptycene-containing PEO-based copolymers, the open clefts of the pentiptycene units may be (partially) occupied by the PEO segments via chain threading, which would impact the free volume fraction and architecture in a way different from other PEO-containing copolymers without iptycene hard segments. Fractional free volume (FFV) of the pentiptycene-PEO-based copolymer films was estimated based on the group contribution method coupled with density measurements, and the results are listed in Table 2. As shown, the FFV values of the copolymer films are in the range of 7.4-12.7% depending on the PEO content, which are much lower than that of the pure pentiptycene–6FDA polyimide (19.7%) without PEO.27 For example, the FFV of pent-PI-PEO2000 (20%) copolymer dropped by 35% of that of the pentiptycene–6FDA polyimide; when the PEO content was increased to 17 ACS Paragon Plus Environment

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60%, FFV of pent-PI-PEO2000 copolymer decreased by 56%. This observation seems to support our previously proposed partial filling mechanism, in which the flexible PEO segments may partially fill both the inter-chain spacing and the internal molecular cavities of pentiptycene units, resulting in significant decreases in FFV.26-27 For the pent-PI-PEO2000 series copolymers, FFV of the PEO (60%) copolymer increases slightly when compared with the PEO (50%) copolymer. This finding is possibly because the dominant rubbery PEO segments increase overall inter-chain spacing, which is discussed in detail later. Given the same PEO weight content, copolymer films with longer PEO sequences showed higher FFV. This effect can be attributed to the higher molar ratio of high free volume pentiptycene units when a higher molecular weight PEO was used for the copolymerization. Table 2. Contact angle, mechanical properties, density and FFV of pent-PI-PEO copolymer films. Polymer

Contact angle

Young’s modulus (GPa)

Elongation at break (%)

Tensile strength (MPa)

Density (g⋅cm−3)

FFV (%)


87

2.60

1.3

35

1.397

11.4


84

2.09

3.0

63

1.380

12.2


86

2.02

2.0

40

1.368

12.7


75

1.63

10

62

1.392

7.9


75

1.45

71

51

1.381

7.4


74

1.33

47

45

1.370

8.7


69

0.35

286

27

1.356

7.7


63

0.33

912

7.4

1.306

8.7

Wide-angle X-ray diffraction (WAXD) measurements were carried out to investigate the chain packing in the copolymer films. Representative WAXD patterns are shown in Figure 6 to demonstrate how chain packing is influenced by PEO content (Figure 6a) and PEO chain length (Figure 6b). In all spectra, no crystalline peaks were observed, especially at 2θ = 19.3° and 23.4°, which are characteristic diffraction peaks of PEO crystals46, confirming the completely amorphous structure of the copolymers. This result is consistent with the DSC results, which showed no melting/crystallization peaks during 18 ACS Paragon Plus Environment


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heating/cooling cycles. Specifically, three amorphous halos located at 2θ = 9.5°, 17.0° and 23.3° can be identified in the spectra, which represent d-spacing values of 9.3, 5.2 and 3.8 Å, respectively. The largest d-spacing value of 9.3 Å is most likely related to the average inter-chain distance induced by the bulky pentiptycene units since this d-spacing value is very comparable with the spatial dimension of the pentiptycene unit.27 This d-spacing value is much larger than most of the reported non-iptycene-containing aromatic polyimides,47-50 indicating that introducing pentiptycene units into the polymer backbone effectively disrupts chain packing leading to high fractional free volume. As shown in Figure 6(a), the relative intensity of the 9.3° peak decreases with increasing PEO content. This is reasonable because there are fewer pentiptycene units in the copolymer main chain structure as PEO weight percentage increases, and at the same time, more clefts of pentiptycene units might be filled at higher PEO contents. Similarly shown in Figure 6(b), the intensity of the 9.3° peak increases with increasing PEO chain length given the same PEO weight content (e.g., 40%) because more pentiptycene units are present in the copolymers with longer PEO sequences. The other two diffraction peaks located at 17.0° and 23.3° with respective d-spacing values of 5.2 Å and 3.8 Å may be assigned to the average inter-chain distance associated with the chain segments without pentiptycene units and the π-π stacking of arene rings of pentiptycene units, respectively.51 The relative intensities of these two peaks are sensitive to PEO weight content, but not PEO chain length. Specifically, the relative intensity of the π-π stacking peak at 23.3° increases with increasing PEO content (Figure 6a). This is because the copolymer backbone becomes increasingly flexible at high PEO content, so the copolymer may assume certain chain conformations that favor the contact of arene rings promoting attractive π-π stacking interactions.


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Figure 6. WAXD spectra of (a) pent-PI-PEO2000 copolymer films with varying PEO content (0% sample refers to the pentiptycene-6FDA polyimide reported in ref. 25); (b) pent-PI-PEO (40%) copolymer films with varying PEO chain length. Gas Permeation Properties. Pure gas permeabilities of PEO-containing copolymer membranes with five gases (H2, N2, O2, CH4 and CO2) were measured using the constant-volume variable-pressure method.34 Permeabilities were measured at upstream pressures ranging from 3 to 17 atm to examine the feed pressure dependence of permeability and explore possible plasticization behavior. As shown in Figure S2 in the supporting information, permeabilities of H2, N2, O2, and CH4 show very weak dependence on feed pressure; while for CO2, especially when the PEO content is high, the permeability increased by approximately 50% when the feed pressure was increased from 3 atm to 17 atm. This trend is a typical sign of plasticization for PEO-containing copolymers with high PEO content.21, 52-53 In the following discussion, we focus on the pure gas transport properties measured at 3 atm and 35 oC. A summary of gas permeability and CO2-related ideal selectivity values for all the copolymer membranes are recorded in Table 3 along with the data of a few comparable PEO-rich copolymers reported previously. Table 3. Pure gas permeability and ideal selectivity of pent-PI-PEO copolymers (3 atm, 35 oC) compared with amorphous pure PEO, pentiptycene-6FDA polyimide and previously reported PEO-rich copolymers.



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permeability (P, Barrer)

ideal selectivity (α)

polymer

a

H2

CO2

O2

N2

CH4

CO2/H2

CO2/N2

CO2/CH4


12

6.6

1.3

0.22

0.19a

0.6

30

—


16

9.0

1.8

0.42

0.36

0.6

21

25


22

19

3.1

0.65

0.76

0.9

29

25


3.6

2.0

0.30

0.05

0.07

0.6

40

29


4.7

4.0

0.49

0.10

0.13a

0.9

40

—


5.3

6.1

0.68

0.16

0.26

1.2

38

23


6.7

16

1.3

0.45

0.72

2.4

36

22


9.6

39

2.7

0.84

2.0

4.1

46

20

pentiptycene-6FDA polyimide27

185

146

31

7.1

5.9

0.8

21

25

amorphous pure PEO7

21

143

8.1

3.0

7.1

7

48

20

6F-ODA/PEO3(80)43

—

19.2

—

0.3

—

—

65

—

6FDA-ODA-PEO1(60)12

9.9

49.4

—

—

—

5.0

—

—

PMDA-mPD/PEO3(80)44

—

99

—

1.98

—

—

50

—

Permeability measured at 17 atm, 35 °C. Gas permeabilities of copolymer films with 20% PEO content (except pent-PI-PEO2000 (20%) film)

follow the order of: P(H2) > P(CO2) > P(O2) > P(N2) > P(CH4), which is consistent with the order of gas kinetic diameters.54 This trend is typical for glassy polymers like these copolymers with low PEO content, which largely rely on molecular size sieving for gas transport, and diffusivity plays a dominant role. However, the sequence of gas permeabilities changes when the PEO content reaches 40 wt% with long PEO2000 sequence. In particular, for CO2/H2 and CO2/N2, gas permeation became more selective towards CO2 with increasing PEO content. For example, pent-PI-PEO2000 (60%) exhibits the highest CO2 permeability of 39 Barrer and the highest CO2/H2 and CO2/N2 selectivities of 4.1 and 46, respectively among all of the copolymer membranes, suggesting favorable contributions of PEO segments towards CO2 related separations. However, an opposite trend was observed for CO2/CH4 21 ACS Paragon Plus Environment

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separation in that CO2/CH4 selectivity decreased slightly with increasing PEO content. It has been demonstrated that high CO2/CH4 selectivities are typically obtained with materials having good diffusion selectivity,55-57 and the incorporation of rubbery PEO segments weakens the diffusion selectivity. As such, the increase in CO2 solubility due to favorable interactions between PEO and CO2 does not make up for the increase in CH4 diffusivity upon addition of rubbery components, resulting in decreased CO2/CH4 selectivities as a function of PEO content. Gas permeabilities also depend on PEO sequence length as shown in Table 3. For a given PEO weight percentage, gas permeabilities increase with increasing PEO molecular weight for all gases. This observation is consistent with the trend of fractional free volume as a function of PEO molecular weight discussed previously in Table 2, suggesting that diffusion selectivity still plays a role in these CO2-philic PEO-based copolymer films due to the strong size sieving ability of pentiptycene-based hard segments. Gas permeability and CO2-related selectivity as a function of PEO content and PEO chain length are plotted in Figure 7 to elucidate the structure-property correlations for these new copolymers. Previous studies demonstrated a roughly linear relationship between CO2 permeability and PEO content for PEO-containing copolymer membranes due to the CO2-selective sorption contribution of the PEO segments.11-13,

44, 58

However, a nonlinear relationship between CO2 permeability and PEO weight

content was observed for these pent-PI-PEO copolymers. For instance, the CO2 permeability of the pent-PI-PEO2000 series decreases from 19 to 6.1 Barrer as PEO content increases from 20% to 40%, and then it increases when the PEO content is higher than 40% and reaches a maximum of 39 Barrer at 60% PEO content (Figure 7a). This trend can be rationalized by the synergetic effect of diffusivity and solubility contributions of PEO segments towards CO2 permeation. At low PEO content (< 40%), partial filling of pentiptycene molecular cavities by PEO segments reduces the diffusive pathway in the copolymer membranes as evidenced by the greatly reduced fractional free volume (Table 2); meanwhile, selective PEO-CO2 interactions are not significant enough to compensate for this loss in diffusivity. Therefore, the copolymers showed reduced CO2 permeability at low PEO content (< 40%). Similar trends were observed for the other four gases. However, when PEO content was greater than 40%, or, more precisely, when the copolymers transitioned from being glassy to rubbery due to 22 ACS Paragon Plus Environment


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incorporation of more PEO sequences, the favorable PEO-CO2 interactions started to contribute appreciably towards increasing CO2 permeability, which is in addition to the increased diffusivity due to flexible chains in rubbery materials. For other non-interacting gases, the permeability increase (at a less profound level) was mainly due to the diffusivity increase when the copolymers became rubbery. As a result, a more profound increase in CO2 permeability was observed when compared with the other four gases because of the combined contributions from both diffusivity and solubility. Accordingly, the ideal CO2/H2 and CO2/N2 selectivities increase almost monotonically with PEO content due to the same synergistic effect (Figure 7b). In particular, the copolymer membranes became reverse-selective in CO2/H2 separation (i.e., α(CO2/H2) > 1, circles in Figure 7b) starting at 40% PEO content with 2000 g⋅mol−1 PEO chain length, which signifies the synergistic contribution of diffusivity and solubility. The effects of PEO chain length on CO2 permeability and ideal selectivities (CO2/H2 and CO2/N2) are illustrated in Figure 7c and 7d, respectively. In general, copolymers with longer PEO chains are more permeable towards CO2 because of higher FFV values (Table 2) in the copolymers when a longer PEO sequence was used. For the dependence of ideal selectivity on PEO chain length, long PEO sequence length is beneficial for CO2/H2 selectivity (Figure 7d). For example, for the same 40% PEO content, the CO2/H2 selectivity of pent-PI-PEO2000 doubles that of pent-PI-PEO600. On the other hand, the CO2/N2 ideal selectivity seems to be less sensitive to PEO sequence length, which maintains similar values regardless of the change in PEO chain length. Consequently, copolymers that possess the highest PEO content and the longest PEO sequence deliver the best separation performance for CO2-related separations.


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Figure 7. (a) Gas permeabilities and (b) ideal selectivities (CO2/H2 and CO2/N2) as a function of PEO weight content for pent-PI-PEO2000 copolymer series; (c) CO2 permeability and (d) ideal selectivities (CO2/H2 and CO2/N2) as a function of PEO sequence length for pent-PI-PEO (40%) copolymers. The gas permeability-selectivity tradeoff plots for CO2/H2 and CO2/N2 gas pairs of the PEO-containing copolymers (35 oC, 3 atm) are presented in Figure 8.59-60 The pent-PI-PEO2000 (60%) film shows the highest combinations of permeability and selectivity for both CO2/H2 and CO2/N2, with a CO2 permeability of 39 Barrer, and CO2/H2 and CO2/N2 selectivities of 4.1 and 46, respectively. For CO2/H2 separation, the copolymer films start to show reverse-selective behavior when the PEO molecular weight is more than 900 g⋅mol−1 and its weight content exceeds 40%, suggesting the importance of both PEO sequence length and weight content in fabricating CO2 selective membranes. It is also important to note that two of the reverse-selective films (pent-PI-PEO900-40% and pent-PI-PEO2000-40%) are still glassy even with 40% PEO, which signifies the supramolecular chain 24 ACS Paragon Plus Environment


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threading interactions in these copolymers that do not exist in other PEO-containing copolymers reported in the literature.12, 14, 35-36 Permeation data of pentiptycene-based polyimide without PEO27 and amorphous pure PEO7 were also included for comparison. For CO2/H2 separation, all the PEO2000-containing copolymer films outperform the non PEO-containing polyimide film in that they all have higher permselectivities. In particular, the pent-PI-PEO2000-60% sample shows both high permeability and high selectivity that approach the CO2/H2 upper bound and the pure PEO membrane. The enhancement in CO2 permeability with PEO incorporation is less pronounced for the CO2/N2 gas pair. However, all copolymer films show higher CO2/N2 selectivities than the glassy pentiptycene-polyimide regardless of PEO chain length and weight content. Additionally, reverse-selective membranes usually show even higher selectivity in mixed gas measurements than single component permselectivity and perform better at low temperatures as demonstrated in several previous studies.12, 21, 61 Further studies on the temperature dependence of transport and mixed gas permeation tests are underway to fully explore the potential of these new CO2-philic membranes.

Figure 8. Permeability-selectivity Upper Bound plots (35 oC, 3 atm) of pent-PI-PEO600 (black squares), pent-PI-PEO900 (red circles) and pent-PI-PEO2000 (blue triangles) copolymers for (a) CO2/H2 and (b) CO2/N2 gas pairs. Arrows in the plots guide the trend of increasing PEO content from 20% to 40% for pent-PI-PEO600 and pent-PI-PEO900 polymers, and 20% to 60% for pent-PI-PEO2000 series. Data points of pentiptycene-6FDA glassy polyimide without PEO (magenta star)27 and amorphous pure PEO (green diamond)7 are included for comparison. 25 ACS Paragon Plus Environment

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4. Conclusions Highly CO2-selective multicomponent copolymer membranes were synthesized from commercial Jeffamine® polyetheramines, pentiptycene-based diamine and 6FDA. The copolymers showed excellent solution-processability

and

were

solution-cast

into

dense

thin

films.

The

resulting

pentiptycene-containing PEO-copolymer films exhibited enhanced mechanical properties as compared with previously reported PEO-copolymers without iptycene structures, which is attributed to the unique supramolecular chain threading and interlocking interactions induced by the pentiptycene units. Partially due to the same effect, the multicomponent copolymer membranes are completely amorphous without noticeable phase separation as revealed by DSC, WAXD and AFM. The fractional free volume shows a strong dependence on both the PEO weight content and PEO chain length. In pure gas permeation tests, the nonlinear trend of CO2 permeability with PEO weight content is ascribed to the competition between solubility and diffusivity contributions, whereby the copolymers change from being diffusivity–selective to solubility–selective when PEO content reaches 40%. Given the same PEO weight content, increased PEO sequence length is beneficial in enhancing both CO2 permeability and selectivity. In particular, CO2 pure gas permeability of 39 Barrer with CO2/H2 selectivity of 4.1 and CO2/N2 of 46 has been obtained for pent-PI-PEO2000 (60%) film, suggesting their potential for CO2-related separations.

ASSOCIATED CONTENT Supporting Information Contact angle of the copolymer films and the feed pressure dependence of pure gas permeabilities.

AUTHOR INFORMATION Corresponding Author *Telephone: +1-574-631-3453 Fax: +1-574-631-8366 E-mail: [email protected]

ACKNOWLEDGEMENTS Ruilan Guo gratefully acknowledges the financial support of the Division of Chemical Sciences, 26 ACS Paragon Plus Environment


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Biosciences, and Geosciences, Office of Basic Energy Sciences of the U.S. Department of Energy (DOE) under Award DE-SC0010330. Shuangjiang Luo would like to thank the partial financial support from the Center of Sustainable Energy at Notre Dame via the ND Energy Postdoctoral Fellowship Program. We thank Haifeng Gao and Yi Shi for the use of SEC, Yingxi Zhu and Benxin Jing for the use of contact angle goniometer, and Chongzhen Na and Haitao Wang for the use of AFM.


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