The evolution of free volume and gas transport properties for the thermal rearrangement of poly(hydroxyamide-co-amide)s membranes

Author’s Accepted Manuscript The evolution of free volume and gas transport properties for the thermal rearrangement of poly(hydroxyamide-co-amide)s membranes Lu Ye, Lina Wang, Xingming Jie, Congyao Yu, Guodong Kang, Yiming Cao www.elsevier.com/locate/memsci

PII: DOI: Reference:

S0376-7388(18)31980-X https://doi.org/10.1016/j.memsci.2018.11.029 MEMSCI16631

To appear in: Journal of Membrane Science Received date: 19 July 2018 Revised date: 2 November 2018 Accepted date: 13 November 2018 Cite this article as: Lu Ye, Lina Wang, Xingming Jie, Congyao Yu, Guodong Kang and Yiming Cao, The evolution of free volume and gas transport properties for the thermal rearrangement of poly(hydroxyamide-co-amide)s membranes, Journal of Membrane Science, https://doi.org/10.1016/j.memsci.2018.11.029 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

The evolution of free volume and gas transport properties for the thermal rearrangement of poly(hydroxyamide-co-amide)s membranes Lu Yea,b , Lina Wanga,*, Xingming Jiea, Congyao Yua,b, Guodong Kanga, Yiming Caoa,* a

Dalian National Laboratory for Clean Energy (DNL), Dalian Institute of Chemical

Physics, Chinese Academy of Sciences, Dalian 116023, China b

University of Chinese Academy of Sciences, Beijing 100049, China

[email protected] [email protected] *

Corresponding authors at: Dalian National Laboratory for Clean Energy (DNL),

Dalian Institute of Chemical Physics, Chinese Academy of Sciences, Dalian 116023, China. Dedicated to the 70th anniversary of Dalian Institute of Chemical Physics, CAS.

Abstract:

Within a microporous polymer membrane, its high gas separation performance is much dependent on the free volume element architecture. In

this

study,

thermally

rearranged

poly(benzoxazole-co-amide)

(TR-PBOA) copolymer membranes were prepared by in-situ thermal treating poly(o-hydroxyamide-co-amide) (PHAA) precursors, basing on commercially available TR-able and non TR-able diamines with different 1

molar ratio. Free-volume topologies were tailored by controlling the degree of thermal rearrangement and the flexibility of the original chains. Upon thermal conversion, small cavities coalesced into bigger ones, representing hourglass-shaped cavities with larger cavities and small bottlenecks, resulting in the significant increase in permeability. It was found that thermal rearrangement mainly occurred near or above glass transition temperature (Tg) where chain segments obtained enough motion ability, and TR-PBOA membrane prepared at this temperature possessed the maximal selectivity due to effective packing of rigid chains. When thermally treated at temperature much higher than Tg, there was a compromise between thermal conversion and chain annealing. Compared to thermal treatment temperature, the effect of dwelling time on thermal conversion ratio was minor, as the formed rigid structure limited chain motion until enough energy was received at higher temperature. Furthermore, TR-PBOA membranes with appropriate ratio of PBO and PA contents displayed superior mechanical properties and gas transport performance, especially for CO2/CH4 separation (CO2 permeability was about 237 Barrer, CO2/CH4 ideal selecitivity was 36.6, plasticization pressure of CO2 was 2.9 MPa) (1 Barrer=10-10 cm3 (STP) cm cm-2 s-1 cmHg-1 ) Keywords：TR-PBOA membranes; thermal conversion temperature; free volume; permeability and selectivity; resistance to plasticization 2

1. Introduction

Recent years, as global energy shortage and other environmental issues become more serious, advanced energy-saving, clean separation and purification technologies are urgently needed. Membrane gas separation process has received considerable attention in air separation, carbon capture, natural gas sweetening as well as hydrogen recovery due to its advantages of low energy cost, small footprint and flexible operation[1] [2] [3]. Microporous polymers have been of great interest for gas separation membranes with their extremely high porosity and easy processibility. PTMSP is considered as the most permeable microporous polymer with CO2 permeablility of 20000-30000 Barrer [4]. This superior permeability derives from its very high free fractional volume (FFV), more than 20% of the total polymer occupied volume, but its ability to separate small gas molecules is too low to be used. Furthermore, the large cavities are susceptible to collapse due to high pressure compression or physical aging. Polymers of intrinsic microporosity (PIMs) and thermally rearranged (TR) polymers are alternative promising microporous materials for separation membranes, whose substantially improved gas separation performance has made great contribution to redefine Robeson’s upper bound in 2008 [5, 6]. PIMs exhibit highly 3

interconnected, irregularly shaped free volumes caused by the chain bending and twisting within rigid macromolecular architecture or robust network [7]. In spite of the exceptional gas permeability, selectivity and durability [8-10], unfortunately PIM membrane preparation is exposed to the problems of monomer impurity and high-cost, multistep synthesis as well as poor solubility, which hinder the application in industry [11]. Thermally rearranged polymer membranes are acquired by in-situ thermal conversion of soluble precursors, and the stiffness and contortion in rigid backbones endows TR polymer membranes with well-tuned cavity size distribution and extraordinary gas separation performance [12, 13]. (a)

(b)

Figure 1 Thermal rearrangement mechanism of (a)TR-α polymer and (b) TR-β PBO[14]

4

The precursors of TR polymers could be polyimides or polyamides containing ortho-functional groups [14]. TR polymers from polyimdes are called TR-α polymers. Depending on the ortho- functional groups to the imide groups, the resultant TR-α polymers can be in the form of polybenzoxazoles (PBO), polybenzothiazoles (PBZ), polypyrrolone (PPL), and polybenzimidazoles (PBI), as showed in Figure 1. The most studied is the thermal conversion of hydroxyl polyimide (HPI) to TR-PBO in the solid state, which occurs over a temperature range of 300–450℃ under inert atmosphere. Upon thermal rearrangement, the formation of rigid rod-like benzoxazole-phenylene structure causes disruption in stiff chains, resulting in the unique bimodal cavity distribution, where the larger cavities provide rapid gas diffusion channels, while the smaller ones are appropriate for molecular sieving [15-17]. Thus, TR-PBO membranes represent significantly enhanced gas permeability and maintained selectivity. It is assumed that cavity size and distribution plays an important role in the separation performance of TR-polymer membranes[18]. It was reported that thermal conversion occurring near or below glass transition temperature (Tg) generates desired microporosity for gas separation [19, 20]. And at higher heating treatment temperature, smaller cavities coalesce into larger ones, forming hourglass shaped cavities with narrow neck regions suitable for separation [15]. The greatest advantage of TR polymer membranes is that 5

their performance can be easily tuned by adjusting templating molecules and heat treatment conditions [14]. Han, et al. investigated the effect of precursor synthesis routes on gas transport performance of TR-PBO membranes and found that tPBO demonstrated ultrahigh CO2 permeability of about 4000 Barrer and CO2/CH4 selectivity of 28 due to the cross-linking of tHPI, surpassing the 2008 polymeric upper bounds for gas separation membranes, whereas aHPI retained a linear chain structure during azeotropic imidization, resulting in aPBO with CO2 permeability of 400 Barrer and CO2/CH4 selectivity of 34 [17]. It was reported ortho-positioned functional groups in polyimide precursors could also influence the thermal conversion temperature and free volume [21]. In spite of the excellent gas transport performance, TR-PBO membranes derived from HPI precursors were often brittle and fragile due to the high treatment temperature, which greatly hindered their application. Various approaches such as cross-linking [22, 23], incorporation of flexible groups [24], copolymerization [25, 26] and fabrication of asymmetric/composite membranes [27, 28] have been employed to improve their mechanical properties. Lately, defect-free and ultrathin TR-PBOI membranes were fabricated by spin-coating a precursor solution of graphene oxide–polyimide (GO-PI) on a porous anodic

aluminum

oxide

substrate

and

subsequently

thermal

rearrangement. Due to the formation of reduce GO scaffold inside TR 6

polymer, the resultant TFC (thin film composite) membrane exhibited enhanced mechanical robustness and remarkable gas permeance without loss of selectivity.[27] Another effective solution to alleviate the problem of fragility and brittleness of TR polymer membranes is to reduce thermal treatment temperature. TR polymer treated at lower temperature can obtain better mechanical properties. In the case of TR-β polymers derived from polyamides, as showed in Figure 1(b), the conversion process could occur at lower temperature of 250~350℃, with no dependence of inert purge condition [29]. And at same thermal conversion temperature, TR-PBO membranes derived from poly(hydroxyamide) (PHA) were more permeable than those from the analogous polyimide precursors due to the higher conversion ratio of PHA. When the thermal treatment temperature was above 400℃，the permeability of TR-α-PBO membranes improved significantly. But the significant improvement in gas transport properties of TR-α-PBO films was usually at the expense of a severe loss in mechanical integrity, whereas TR-β-PBO films still maintained the qualitative mechanical flexibility and ductility [30]. Moreover, the synthesis process of polyamides is simpler than that of polyimdes. It reveals that the PHA cyclodehydration route is superior in terms of synthesis cost, manufacturing technics and mechanical properties. Deep study in PHA-to-PBO route would be of great sense to boost gas transport 7

performance and hollow fiber membrane preparation for industrial application. In this study, we investigated the evolution of free volume elements and gas transport properties during the thermal conversion from PHA to PBO. Here, the precursors of poly(hydroxylamide) were synthesized from commercially available monomers: 2,2'-bis(3-amino-4-hydroxyphenyl) hexafluroprop (BisAPAF) and terephthaloyl chloride (TPC). Also, 4,4'-diaminodiphenyl ether (ODA) containing flexible ether group was used for copolymerization to regulate the mechanical and gas separation properties of TR-polymer membranes, and also explore the effect of non-TR-able

segments

on

thermal

conversion

process.

Poly(hydroxyamide-co-amide) (PHAA) precursor films were thermally treated at different temperature, thus a set of TR-polymer membranes with varying PBO content were obtained representing various stage of thermally rearranged process. ATR-FTIR of all films samples were characterized to verify the structure change; WAXD was measured to explore the pore distribution and evolution during thermal conversion process; DMA and TG were conducted to investigate the relationship between Tg and thermal cyclization temperature, as well as changes in thermal properties; the thermal rearranged behavior of PHAA samples at certain temperature was also monitored by TG analysis; densities and XPS were measured to calculated FFV; tensile testing was conducted to 8

character the mechanical properties; gas permeability and plasticization pressure were obtained using constant-volume / variable-pressure method.

2. Experimental

2.1. Materials Terephthaloyl

chloride

(TPC)

(99%),

2,2'-bis(3-amino-4-hydroxyphenyl) hexafluoropropane (BisAPAF) (98%) and 4,4'-diaminodiphenyl ether (ODA) (98%) were products of Tokyo Chemical Industry Co., Ltd (Japan). Prior to synthesis, bisAPAF and ODA were dried in vacuum oven at 120℃ and 40℃ respectively for 24h to remove the absorbed water and impurity substance. Anhydrous N-methyl-2-pyrrolidone (NMP) (analytical grade) was supplied by Kermel Chemical Reagent Co., Ltd (Tianjin, China), which was treated with molecular sieves before use. 2.2. Synthesis of copolymer precursors The

precursor,

poly(hydroxyamide-co-amide)

(PHAA),

was

synthesized from TPC, bisAPAF and ODA via low temperature polycondensation routine, as showed in Figure 2.

9

Figure 2 The synthesis scheme of poly(hydroxyamide-co-amide)s (PHAAs) and their thermally rearranged poly(benzoxazole-co-amide)s (TR-PBOAs)

In a 100-mL three-necked round bottom flask, 10 mmol of diamines with different molar radio of bisAPAF and ODA were added and stirred in 18 ml NMP under N2 atmosphere. The molar ratio of bisAPAF and ODA for copolymerization was adjusted to 10:0, 8:2, 5:5, 2:8 and 0:10 respectively. Then, the completely dissolved solution was cooled to 0℃ in an ice bath and a stoichiometric amount of TPC was added. As the viscosity of the solution increased with reaction time, the flask was then 10

placed in oil bath and heated to 80℃ to increase the activity of unreacted monomers. After holding for 4h at 80℃, the obtained brownish viscous solution was cooled to room temperature and precipitated in 3:1 (v/v) deionized water-ethanol solvent. The precipitate was then collected and washed several times with deionized water and finally dried in a vacuum oven at 80℃ for 24h. The solubility of the PHAAs was qualitatively tested and the result was summarized in Table 1. Due to their poor solubility in most organic solvents, PHAA with high ODA content (2:8 and 0:10) could not be fabricated into effective membranes. 2.3. Preparation of dense membranes Precursor solutions, 6 wt% concentration in NMP, were filtered using filter cloth (100 meshes) to remove undissolved materials and dust particles. Then the degassed solutions were cast on leveled clean glass plates at 60℃ for 24h to remove the bulk of solvent. Subsequently, the films with a diameter of 8.0 cm were peeled off from glass plates and dried in a vacuum oven at 80℃ for 12 h, 120℃ for 4h and 160℃ for 24h to remove residual solvent. The obtained PHAA films were heated to target temperature at 5℃/min and maintained for 1h in muffle furnace with N2 purging. The resultant thermal rearranged polymer membranes were designated as TR X-Y, where X refers to thermal treatment 11

temperature, and Y refers to duration of thermal treatment. For example, TR 350-1(5:5) signified that PHAA (5:5) film experienced thermal treatment of 350℃ with a duration of 1h. 2.4. Characterization of physical property Attenuated total reflectance Fourier transform infrared (ATR FT-IR) spectra was characterized by a Thermo Scientific Nicolet iS5 Fourier transform infrared spectrometer (Waltham, Massachusetts, USA). Wide angle X-ray diffraction (WAXD) was conducted on X’pert Pro-1 X-ray diffractometer (PANalytical, the Netherlands) in reflection mode, using CuKα radiation at a wavelength 1.54 Å, operating in a 2θ range of 5–40° with a scan rate of 5° min-1. Dynamic mechanical analysis (DMA) was carried on the TA Instruments-TA Q800 to obtain glass-rubber transition temperature (Tg) and modulus. All samples were measured in temperature sweep mode (5℃min-1 from room temperature to 390℃, 1Hz) under N2 atmosphere. Thermogravimetric analysis (TGA) was performed on a NETZSCH STA 449F3 thermal gravimetric analyzer. PHAA and PBOA samples were conducted with a rate of 10℃min-1 from 40℃ to 900℃ to confirm the thermal conversion conditions and the thermal stability. PHAA (5:5) was also performed at 250℃ and 350℃ for 3h on the TG instrument respectively to monitor the thermal treatment process. 12

In most studies [13, 31], the degree of TR conversion was obtained according to TGA graph of precursors. However, considering that the actual mass loss was much higher than the theoretical mass loss, X-ray photoelectron

spectroscopy

(XPS)

analysis

was

conducted

by

Thermofisher SCALAB 250Xi (USA) to characterize the chemical binding energy levels of C, N, O, F. And TR conversion ratio w was calculated by the following equation: w=

AC=N

1

AC=N +ACO−NH +ANH2

where A represented the peak area of a specific functional group’s binding energy in TR-PBOI samples. Fractional free volume (FFV) was calculated from the equation: FFV =

𝑉−𝑉0 𝑉

=

𝑉−1.3𝑉𝑤

2

𝑉

where, the specific volume of polymer V (cm3 g-1) was measured by the density gradient method. The van der Waals volume (cm3 g-1) was obtained from Bondi’s group contribution method [2]. In this study, irrespective of possible degradation or other reactions at higher temperature, the partly converted PBOA contained three parts: 1) thermally rearranged PBO segment, 2) unconverted poly(hydroxyl amide) segment

(bisAPAF-TPC);

3)

non-TR-able

polyamide

segment

(ODA-TPC). So the effective occupied volume of PBOA, V0 (cm3 g-1), was the sum of the occupied volumes of these three parts [32, 33]: V0 = x[wV0,t + (1 − w)V0,u ] + (1 − x)V0,n

3 13

where, V0,t was the characteristic occupied volume of TR PBO-based polymer, V0,u was the characteristic occupied volume of unconverted PHA-based polymer (bisAPAF-TPC), V0,n was the characteristic occupied volume of non-TR-able PA-based polymer (ODA-TPC), x was the molar fraction of bisAPAF in total diamines, and w was the degree of TR conversion calculated by equation 1. Mechanical properties were measured using XLW Auto Tensile Tester (Labthink®, Jinan, China). Samples for tensile testing were cut into strips of 8mm×80mm, with thickness between 50-80μm. After a sample was mounted, it was pulled at 100 mm/min until broken. Then tensile stress and elongation at break were obtained. 2.5. Gas permeability measurement The gas transport properties were measured using variable–pressure constant-volume method [34]. The gas permeability of five representative gases H2, N2, O2, CH4, CO2 (>99.99% purity) was calculated by the following equation: P=

dp

T VL

0 ( ) dt P TAP 0

4

where P is the gas permeability coefficient, in a unit of Barrer (1 Barrer=10-10 cm3 (STP) cm cm-2 s-1 cmHg-1), dp/dt is the rate of pressure rise under the steady state, V (cm3) is the downstream volume, L (cm) is the membrane thickness, T (K) is the measurement temperature, A (cm2) 14

is the membrane area, P (cmHg) is the upstream pressure, and T0 (K) and P0 (cmHg) are the standard temperature and pressure. As mentioned in Equation 1, permeability is the product of solubility and diffusivity. The diffusion coefficient D (cm2 s-1) was derived from the time lag method, represented by the equation of D=

L2

5

6θ

where θ is the time lag. Then the solubility coefficient S (cm3 (STP) cm-3 cmHg-1) was acquired by S=P/D. Ideal selectivity α of gas pairs was the ratio of permeabilities of two 𝐷

gases, which could be divided into two parts: diffusion selectivity ( 1 ) 𝐷2

𝑆

and solubility selectivity ( 1 ), as showed in the following: 𝑆2

α=

P1 P2

=

D1 D2

×

S1 S2

6

For the investigation of resistance to plasticization, the CO2 feed pressure was increased incrementally, and the permeability was measured at each successive increasing pressure increment.

3. Results and discussion

3.1. Synthesis of PHAAs Different PHAAs were obtained from diamines with different molar ratio. The precipitated PHAA (10:0) as 2-3cm short fibers was soluble in most organic solvents, while PHAA (2:8) and (0:10) in form of robust and 15

consecutive fibers demonstrated poor solubility, as indicated in Table 1. It revealed the tightly alignment of ODA-TPC part prevented the infiltration of solvent molecules, leading to poor solubility.

Table 1 Solubility behavior of precursor polymers PHAA solibilitya NMP DMF DMAc DMSO Aceton Ethano CHCl3 e l ＋＋ ＋＋ ＋＋ ＋＋ －－ －－ －－ 10:0 ＋＋ ＋＋ ＋＋ ＋＋ －－ －－ －－ 8:2 ＋＋ ＋＋ ＋＋ ＋＋ －－ －－ －－ 5:5 ＋－ －－ －－ －－ －－ －－ －－ 2:8 ＋－ －－ －－ －－ －－ －－ －－ 0:10 ＋＋: soluble at room temperature;＋－: swelling or partially dissolving; insoluble even upon heating; a: solubility was tested with 10 mg of precursor samples in 5 mL solvent

THF －－ －－ －－ －－ －－ －－:

3.2. Structural characteristics ATR-FTIR measurements were performed on the precursors of PHAAs and their thermally treated samples at different temperature to follow the evolution of structural change during the copolymerization and thermal rearrangement, as shown in Figure 3. PHAA-to-PBOA conversion was demonstrated by the disappearance of O-H vibration in the broad region of 3200-3400 cm-1, the attenuation of characteristic adsorption bands of amide [1644 cm-1 (-C=O) and 1500 cm-1(-C-N)] in precursors, as well as the appearance of benzoxazole ring-structure characteristic bands at wave number 1558 cm-1 (-C=N), 1477 cm-1 (-C=N), 1058 cm-1 (-C-O) and 921 cm-1 [35]. When raising thermal treatment temperature, the content of 16

PBO moiety in backbones increased with thermal conversion ratio, so the adsorption intensity of benzoxazole bands became stronger, while that of amide bands decreased or vanished. For PHAA (10:0), the characteristic amide bands were still present at 300℃ and became imperceptible upon 350℃ treatment, indicating partial conversion of precursors at 300℃. Given that PHAA (8:2) and (5:5) containing non-rearrangeable building blocks, the characteristic amide bands wouldn’t disappear completely if no degradation took place. All samples treated at 350℃ showed an almost negligible peak at 1500 cm-1 (C-N of amide group), indicating an almost complete conversion at around 350℃. Figure 3 (d) compared the structure constitution of different PHAAs. With an increase of ODA moiety in precursor copolymer, the absorption intensity at 1215 cm-1 (C–O stretching in the ether linkage) became stronger. Besides, C=O and C–N band of amide unit both shifted to lower wave number region (from 1650 cm-1 to 1646 cm-1 and from 1509 cm-1 to 1497 cm-1 respectively) due to inductive effect of the oxygen atom in ODA unit.

17

(a)

(b)

18

(c)

(d)

Figure 3 ATR-FTIR spectra showing structural changes of (a) TR process for PHAA(10:0), (b) TR process for PHAA (8:2), (c) TR process for PHAA (5:5) and (d) precursors from different bisAPAF/ODA ratio.

Figure 4 showed WAXD patterns of precursor films and converted 19

films treated at different temperature. According to Bragg's equation ( d = λ/2sinθ ), the average d-spacing values of the polymers were calculated and listed in Figure 4. The presence of broad and weak amorphous halos in all cases indicated both PHAA and PBOA films were in the amorphous state [36]. The PHAA (10:0) precursor film displayed an amorphous peak at 17.1°, corresponding to a distance of 0.518nm. As the ODA content in PHAA increased, the molecular chains were more flexible and tight-packing, whose intersegmental distance reduced. After thermal treatment, the patterns differed in both shape and width of the halos. Upon 250℃ heat treated, PHAA (10:0) partly converted into TR-PBOA, which exhibited two broad diffraction peaks located at 2θ = ~14.8°

and

~21.0° ,

in

accordance

with

proposed

hourglass-shaped cavity mode. The same case was also observed for (8:2) series and (5:5) series. On the one hand, during thermal rearrangement process, the incorporation of benzoxazole moieties into the polyamide chain backbones formed rigid benzoxazole-phenylene coplanar structure. The formed rigid rod disturbed chain packing, and led to larger interchain distance. On the other hand, thermal conversion reduced the amount of hydroxyl bonds and lessened intermolecular forces among PBO chain units. Thus some atoms of the rod-like chains could get closer, forming narrow bottleneck regions [37]. Also, interchain cross-linking was probably occurred when temperature increased [38], also contributing to 20

the development of two diffraction peaks. However, as thermal treatment temperature further increased up to 350℃, the original two peaks in X-ray diffraction patterns of PHAA (8:2) and (5:5) became inconspicuous, while PHAA (10:0) still maintained its bimodal distribution due to rigid structure and superior thermal stability. As the heat treatment temperature reached 400 ℃ , all samples presented a broad and flat solo peak, indicating a complete disordered state due to chain relaxation and partial degradation at higher temperature. It seemed that thermal treatment temperature had a significant influence on the pore size and distribution. As confirmed from DMA results the following discussion, tanδ peak of glass-rubber transition was observed at lower temperature than that of thermal rearrangement process. Here, referring to previous literature[20, 23, 35], temperature range (200-450℃) of thermal treatment process was divided into three intervals: (1) from TTR1 (the onset temperature of thermal conversion) to Tg (the glassy transition temperature); (2) from Tg to TTR3 (the offset temperature of thermal cyclization); (3) above TTR3. The first interval was related to thermal reaction before chain segments received sufficient mobility, forming more out-of-equilibrium free volume elements (FVEs) derived from chain disruption in stiff chains. Previously studies also demonstrated that thermal conversion occurring near or below Tg could obtain desired micropores [19]. The second stage occurred above Tg where the 21

restriction on segmental motion was removed, thus the relaxation of nonequilibrium chain conformations eliminated larger voids, reducing the discrepancy of microvoid size [7]. As for the thermal treatment above TTR3 where thermal conversion was complete, the effect chain packing was disturbed, resulting in a wide pore distribution.

(a)

22

(b)

(c)

Figure 4 Wide angle X-ray diffraction (WAXD) patterns of (a) (10:0) series, (b) (8:2) series, and (c) (5:5) series.

23

3.3 Thermal properties of PHAAs and TR-PBOAs (a)

(b)

24

(c)

Figure 5 Dynamic mechanical results of PHAAs: (a) tanδ; (b) storage modulus (c) loss modulus as a function as temperature under 1Hz.

Figure 5 displayed the dynamic mechanical results of PHAAs, including the curves of tanδ and modulis as a function of temperature under 1 Hz. The distinction of the three samples could be well explained by their differences in structure construction and molecular mobility [39]. According to Figure 5 (a), only PHAA (10:0) sample had a small peak at 90℃，corresponding to localized sub-glass motions of limited range. Because ODA-TPC moieties in copolymers of PHAA (8:2) and (5:5) acted as anchor points at lower temperature due to the strong molecular interaction and tighter chain packing, which hindered sub-glass motion, PHAA (8:2) and (5:5) had higher storage modulus below 200℃. However, as temperature increased, the storage modulus of PHAA (8:2) and (5:5) 25

declined slightly while that of PHAA (10:0) maintained. It seemed that temperature had greater influence on molecular interaction of chains containing ODA moieties. Due to the presence of flexible ODA, the glass-rubber transition of PHAA (8:2) and (5:5) mainly occurred at 240℃——lower than 270℃ of PHAA(10:0)—— supported by the appearance of tanδ peak as well as the declined storage modulus and increased loss modulus at this temperature. In Figure 5 (a) , the tanδ peaks of PHAA(8:2) at 300℃ and PHAA(5:5) at 315℃ were corresponded to thermal rearrangement process, which was evident by the recovery in modulus due to the formation of rigid structure. For PHAA (10:0), the glass-rubber transition temperature and thermally rearranged temperature were so close that the peaks overlapped each other. Moreover, during the dynamic mechanical heating sweep of PHAA (5:5), a peak at higher temperature of 355℃ appeared along with a strong drop in modulus, which was scribed to the segmental movement of chains in TR polymer. Same trend was also observed in PHAA (8:2) at 380℃. As the highest scan temperature was 390℃, the motion of TR section in PHAA (10:0) was not detected. (a)

26

(b)

(c)

27

(d)

Figure 6 Dynamic mechanical results of PHAAs and PBOAs treated at different temperature: (a) tanδ of (10:0) series; (b) tanδ of (8:2) series; (c) tanδ of (5:5) seires; (d) storage modulus of (10:0) series.

28

Samples of PHAAs treated at different temperature were also measured by DMA analysis, as showed in Figure 6. For TR samples, the sub-glass motion was still evident, while the glass transition process was obviously absent, with no indication of tanδ peaks being observation. And the peaks of thermal conversion vanished, except for TR 250-1 (5:5) and 300-1 (5:5), who exhibited a small amount of thermal rearrangement, in accordance with the higher TR tanδ peak of PHAA (5:5). As the scan temperature reached to 380℃，all samples of TR 250-1 and 300-1 got enough ability of chain motion, illustrated by the rise of tanδ and the drop in storage modulus. TR 350-1 and 400-1 demonstrated fine thermal stability due to their higher benzoxazole contents.

(a)

29

(b)

(c)

(d)

30

Figure 7 Thermogravimetirc analysis for (a) PHAA copolyemrs, (b) (10:0) series, (c) (5:5) series from 40℃ to 900℃, and (d) PHAA (5:5) at certain heating temperature for 3h.

Table 2 Thermal properties of PHAAs Tg (℃)

TTR1 (℃)

TTR2(℃)

TTR3(℃)

Td5%(℃)

PHAA(10:0)a

386

238

333

441

628

PHAA(8:2)

a

320

237

312

402

518

PHAA(5:5)

a

290

232

285

339

453

PHAA(2:8)

a

270

343

440

328

444

Samples

278

230

a

274

275

PHAA(10:0)b

270

215

279

375

523

PHAA(8:2)

b

240

213

277

372

516

PHAA(5:5)

b

240

200

263

368

498

PHAA(2:8)

b

198

260

358

494

PHAA(0:10)

Tg was obtained from tanδ of DMA TTR1 was the onset temperature of thermal cyclization TTR2 was the greatest weight loss temperature during thermal cyclization TTR3 was the offset temperature of thermal cyclization Td5% was 5% weight loss temperature after the fully thermal cyclization in the TGA graph. a: PHAA in reference [35] 31

b: PHAA in this study.

Figure 8 Possible reactions occurred along with the thermal rearrangement of PHAAs.

Thermal stability of a series of PHAA and PBOA membranes with different ODA content was investigated by means of thermogravimetric analysis. According to Figure 7(a), all PHAA precursors had two-stage weight loss TGA profiles. The first weight loss step located in the range of 200-350℃ was corresponded to the thermal rearrangement process where two molecules of H2O were released per repeat unit. The second weight loss step was ascribed to primary thermal degradation of formed PBOAs [40]. It was noted that experimental total weight losses in all cases during thermal conversion process were much higher than the theoretical ones. It has been evidenced by Tai-Shung Chung’s group that a small amount of CO2 was evolved along with the thermal cyclization reaction [38]. They suggested that hemolytic reactions or hydrolytic reactions might take place, as shown in Figure 8. In this study, the 32

maximum weight loss occurred at about 260~280℃, which was a little higher than Tg of PHAAs according to DMA results, inflecting that thermal rearrangement reaction was facilitated by segmental motion. Like other PHA precursors, thermal cyclization of PHAAs occurred at lower temperature than that of HPI polymers due to the linear amide linkage. In contrast, with the tertiary amine in imide linkage and contorted non-planar structure, HPI precursors was demonstrated higher thermal rearranged temperature due to the higher rotational barrier energy of ortho-hydroxyl [30]. It was also found that PHAA with higher ODA content had the lower initial / final thermal conversion temperature, and inferior thermal stability according to the degradation temperature, which was in accordance with the work by Yu Seong Do [35], as shown in Table 2. Figure 7 (b) and (c) showed TGA curves of PHAA and TR-PBOA membranes. It was observed that the char yield at 900℃ increased by 25% after thermal conversion, confirming that the thermal stability was significantly improved upon thermal treatment. However, the initial decomposition temperature did not change significantly for all samples, which increased slightly when raising heat treated temperature from 250℃ to 350℃, but decreased at 400℃. As mentioned above, thermal conversion had been completed at about 370℃, so the loss of pendant groups or partial degradation of backbones might occur, leading to inferior thermal 33

stability. Figure 7 (d) displayed the thermal changes during thermal treatment at certain temperature. With heating treatment at 250℃ for 3h, the weight decreased rapidly within the initial first hour, but then maintained almost constant with heating time, indicating that initial thermal conversion rate was very rapid [41]. The rapid initial conversion was due to the good chain flexibility of PHAA polymer. However, once the rigid benzoxazole structure was formed, the Tg of PBOA was increased and chain motion was limited, preventing the further thermal conversion. While increasing temperature could restart the thermal rearrangement, according to the higher weight loss at 350℃. 3.4. Calculation of FFV As mentioned above, it was hard to determine the thermal conversion ratio by TGA graphs of precursors, due to other reactions occurring along with thermal rearrangement. Instead, XPS was conducted and thermal conversion ratio was calculated following Equation 1. Figure 9 exhibited the N 1s spectra of PBOA samples, the binding energy of -CO-NH- and -CO-NH2 was around 400.1-400.3 eV, and the binding energy of -C=Nwas about 399.2eV for (5:5) series and 399.0 eV for (10:0) series. The calculated thermal conversion ratio and FFV were listed in Table 3. It was noticed that the thermal conversion ratio of 400-1 (10:0) was 55%, much 34

lower than other samples. Because at 400℃, hemolytic and hydrolytic reaction became increasingly important [42].

(a)

(b)

35

(c)

(d)

36

(e)

(f)

37

(g)

(h)

38

Figure 9 XPS N 1s spectra of the surfaces of (a) 250-1 (5:5), (b) 300-1 (5:5), (c) 350-1 (5:5) , (d) 400-1 (5:5), (e) 250-1 (10:0), (f) 300-1 (10:0), (g) 350-1 (10:0) and (h) 400-1 (10:0).

Table 3 Physical properties of PHAAs and TR-PBOA membranes samples

density(g/cm3)

conversion(%)

V(cm3/g)

Vw(cm3/g)

FFV

PHAA(5:5) 250-1(5:5) 300-1(5:5) 350-1(5:5) 400-1(5:5)

1.403 1.386 1.369 1.352 1.339

0 67 89 91 93

0.7125 0.7214 0.7306 0.7396 0.7466

0.4487 0.4529 0.4544 0.4545 0.4547

0.181 0.184 0.191 0.201 0.208

PHAA(10:0) 250-1(10:0) 300-1(10:0) 350-1(10:0) 400-1(10:0)

1.440 1.412 1.406 1.394 1.390

0 65 78 97 55

0.6944 0.7082 0.7112 0.7172 0.7192

0.4257 0.4316 0.4328 0.4347 0.4307

0.203 0.208 0.209 0.212 0.222

39

3.5. Mechanical properites Table 4 Mechanical properties of PHAAs and their thermally rearranged PBOAs samples Tensile stress (MPa) Elongation at break (%) 94 6 PHAA(5:5) 114 6 300-1(5:5) 52 4 400-1(5:5) 56 6 PHAA(10:0) 32 4 300-1(10:0) 0.3 0 400-1(10:0)

Mechanical properties of PHAA and TR-PBOA membranes were presented in Table 4. Tensile stress and elongation at break of (5:5) series were significantly larger than those of (10:0) series, signifying that the ODA made membranes more robust and tough. Compared with PHAA membranes, tensile stress of 300-1 (10:0) was decreased, and even 400-1 (10:0) sample was too fragile to be measured. While that of 300-1(5:5) was improved due to the combination of rigid and flexible structure. It was revealed that mechanical properties could be optimized by tuning thermal conditions and chemical structure. 3.6. Gas transport performance Table 5 Gas transport properties of PHAAs and their thermally rearranged PBOAs at 35℃, 0.2MPa. Permeability（Barrer） Ideal selectivity Samples H2 N2 O2 CH4 CO2 O2/N2 CO2/CH4 CO2/N2 PHAA(5:5) 250-1(5:5) 300-1(5:5) 350-1(5:5) 400-1(5:5) PHAA(8:2)

5.343 0.053 0.321 0.030 1.345 34.4 0.685 3.19 0.37 14.9 62.0 1.24 6.71 0.91 33.8 62.2 1.52 7.61 1.04 36.4 288 10.1 47.6 6.49 237 7.153 0.070 0.403 0.032 1.612

6.09 4.65 5.41 5.01 4.69 5.75

44.23 40.34 37.09 35.04 36.61 50.46

25.51 21.72 27.22 23.97 23.42 23.03 40

250-1(8:2) 300-1(8:2) 350-1(8:2) 400-1(8:2) PHAA(10:0) 250-1(10:0) 300-1(10:0) 350-1(10:0) 400-1(10:0)

76.6 107 148 605 10.4 108 196 187 291

1.90 8.64 1.07 2.49 12.4 1.37 4.28 21.6 2.66 20.0 82.7 15.1 0.095 0.578 0.041 2.71 12.8 1.50 5.82 27.9 3.51 5.71 26.7 4.01 11.6 47.7 7.79

38.0 52.8 104 382 2.23 55.8 120 113 225

4.56 4.96 5.06 4.13 6.11 4.71 4.78 4.67 4.13

35.46 38.66 39.07 25.34 54.28 37.32 34.14 28.26 28.91

20.06 21.21 24.32 19.09 23.53 20.59 20.57 19.83 19.49

Pure gas permeabilities of H2, N2, O2, CH4 and CO2 were measured at 35℃ and 0.2 MPa as listed in Table 5. For all PHAA samples, they exhibited extremely low gas permeability due to much denser chain packing induced by strong hydrogen bonding between the ortho-hydroxy and amide groups. After converting to PBOA membranes, the permeability of obtained films increased significantly. Upon 250 ℃ treatment，the gas permeability of PBOA membranes showed one order of magnitude increase for PHAA (5:5), while almost two orders for PHAA (10:0), demonstrating that TR copolymers with high PBO moieties possessed higher gas permeability as more barrier-free pathways were presented. Thermal conversion at 300℃ resulted in an increase in both permeability and selectivity compared to 250-1 series, due to the effective segmental movement and rearrangement according to DMA analysis. On the one hand, thermal rearrangement was mainly occurred at nearly 300℃, thus more free volume elements for gas permeation were formed at this temperature. On the other hand, above Tg, effective segmental movement could eliminate the formed large cavities, benefit to the 41

enhancement in selectivity. Further increasing heat treatment temperature up to 350℃, not much change in permeability and selectivity could be seen, indicating that 300℃ could be the most suitable temperature for thermal treatment. What’s more, PHAA membranes treated at 400℃ exhibited most permeable performance, but their capability to separate gas molecules was far away from satisfactory due to the partial degradation above full conversion temperature. However, the selectivity of TR 400-1 (10:0) was a bit higher than TR 350-1 (10:0) in terms of CO2/CH4 separation, because the selectivity reduction due to degradation might be offset by the selectivity increase resulting from relaxation of polymer chains at higher temperature. (a)

(b) 42

Figure 10 (a) CO2 permeability and ideal selectivity of CO2/CH4 ; (b) O2 permeability and ideal selectivity of O2/N2 for membranes of TR (5:5) series treated at different temperature and dwelling time.

Figure 10 displayed gas transport properties of TR (5:5) series treated at different temperature and dwelling time. As confirmed from Figure 8(d), the formation of rigid benzoxazole structure limited chain motion, preventing further thermal conversion even with prolonged dwelling time; therefore gas permeability of TR-PBOA membranes almost remained unchanged at constant thermal treatment temperature with different dwelling time. But at higher thermal treatment temperature, as chain segmental motion got enough energy, thermal rearrangement restarted accompanied with a jump in permeability between samples of TR 250-3 and TR 350-1. However, thermal treatment temperature only showed a 43

modest effect on selectivity. The respective permeability-selectivity relationship of PHAA and TR-PBOA membranes was compared against Robeson's upper bounds [43] [6]. Figure 11 showed the plots for CO2/CH4 and O2/N2 gas pairs. Compared to PHAA series, the gas transport properties of TR series moved to right in the figure, indicating the significant increase in permeability after thermal conversion. In spite of the trade-off relationship between permeability and selectivity, the permeability of TR polymer membranes greatly increased without substantially reduced selectivity compared to their precursors, many surpassing the 1991 CO2/CH4 separation limits. Especially for TR 400-1(5:5), whose CO2 permeability was about 237 Barrer with the selectivity of 36.6 over methane, closer to the 2008 upper bound compared to other TR-β PBO membranes [35, 38, 40, 44]. (a)

44

10,000

(5:5) series (8:2) series

Upper bound revised (2008) Upper bound (1991)

CO2/CH4

1,000

100

10

PHAAs

1 0.001

0.1

10

1000

100000

P(CO2) /barrer

(b) 100 (5:5) series

Upper bound revised (2008)

(8:2) series

O2/N2

Upper bound (1991) 10

PHA 1 0.001

0.01

0.1

1

10

100

1000 10000 100000

P(O2) /barrer

Figure 11 Gas transport property of PHAA and PBOA membranes compared with 1991 and 2008 upper bounds for gas pairs of (a) CO2/CH4 and (b) O2/N2. (○ represented TR-βpolymers with HAB moiety [35], □ represented other TR-β polymers with bisAPAF moiety [38, 40, 44, 45])

As the transport of gas molecules through polymer membranes was governed by the well-known “solution-diffusion” mechanism, to study the gas transport performance of converted films in depth, diffusivity was 45

derived from the time lag method and solubility was calculated from the direct ratio between permeability and measured diffusivity. (a)

(b)

46

(c)

Figure 12 Diffusion coefficients of (a) (10:0) series, (b) (8:2) series and (c) (5:5) series films against the kinetic diameters of five light gases

47

Diffusion coefficient was plotted against the kinetic diameters of five light gases (H2, N2, O2, CH4 and CO2), as shown in Figure 12. Diffusivity decreased with increasing kinetic diameter of gases (except CO2), indicating that both PHAA and PBOA membranes owned the ability to sieve penetrant molecules based on their size. The unusual CO2 diffusion behavior was due to the strong interaction between CO2 and polar groups, which was commonly seen in glassy polymeric membranes [46]. Upon 250℃ thermal treatment, diffusivity increased dramatically due to the formation of PBO. The disruption of rigid rod-like PBO structure in matrix caused larger FFV, which benefited to increase pathway of gas transport. Further enhancing heat treatment temperature, chain annealing above Tg offset the increased free volume elements derived from thermal rearrangement, so reduced diffusivity was observed. As mentioned above, degradation occurred when further increasing the rearrangement temperature up to 400 C, leading to a significant increase in diffusion along with the formation of macrovoids. It should be recognized that diffusion coefficients were one or two orders of magnitude bigger than those of precursor membranes, while the FFV only increased by 2-11%, as diffusivity was not only dependent on FFV but also related to the size and distribution of free volume elements. The model of cavity coalescence could explain this phenomenon: cavity radius of precursor membranes was much smaller with a broad 48

distribution of cavity size, and some smaller cavities were assumed to coalesce into bigger ones upon thermal treatment, representing hourglass-shaped cavities. The higher diffusivity and permeability of TR-PBOA membranes was due to the higher amount of interconnected free volume with larger cavity size. The narrower bottleneck diameters in free volume elements were benefit to maintain the selectivity. Thus, in spite of similar FFV, significant increase in diffusivity was observed due to the specific shape and distribution of free volume elements.

49

Figure 13 Solubility of (a) (10:0) series, (b) (8:2) series and (c) (5:5) series films against the critical temperatures of five light gases

50

Figure 13 showed the relationship of solubility and critical temperatures of the gases. In general, the solubility of all samples showed a linear relationship with critical temperature, which was the typical phenomenon in most glassy polymeric membranes[47-49]. On the one hand, the increase in solubility after the TR reaction was ascribed to enhanced FFV with the larger sorption capacity. On the other hand, compared to precursors, the interaction between membrane and gas molecules reduced slightly as the thermal rearrangement proceeded, due to the absence of a side chain or polar groups in TR polymers [50]. As a matter of fact, the reduced interaction offset the enhanced sorption capacity after thermal conversion, thus the improvement in solubility was minor, signifying that diffusivity increment had more contribution to the remarkably high permeability of TR-PBOA membranes. Furthermore, decreased solubility selectivity and increased diffusion selectivity was observed after thermal rearrangement. 3.7 Resistance to plasticization Many high performance glassy polymer membranes were susceptible to plasticization at high level of CO2 or hydrocarbons. Plasticization might be interpreted as an effect of reducing the interaction between adjacent segments of neighboring polymer chains and thus the swollen polymer membrane lost its selectivity [51]. It was considered that the 51

onset of plasticization was caused by excessive polymer swelling, which was relate to the sorbed CO2 partial molar volume in polymers. When the sorbed CO2 partial molar volume reached a threshold value, the improved local chain segmental mobility caused changes in size and frequency of transient gaps in polymer matrix, resulting in enlarged diffusion coefficient. The most common approaches for stabilizing these materials were cross-linking and thermal treatment. [52]

Table 6 Plasticization pressures of PHAA and TR-PBOA polymer membranes. Plasticization Testing temperature(℃) samples pressure (Mpa) 1.95 35 PHAA (5:5) 2.2 35 300-1 (5:5) 2.9 35 350-1 (5:5) 2.0 35 PHAA (10:0) 2.3 35 300-1 (10:0) 2.5 35 400-1 (10:0) 1.2 22 Matrimida a 2.2 23 P84 a 2.8 21 Polyetherimide a 2.4 23 BPZ-polycarbonate a 3.1 25 BPA-polycarbonate a 3.4 23 Polysulfone a: the plasticization pressures were referred to reference[53].

The effect of plasticizing agent on the permeability of PHAA and TR-PBOA membranes was investigated using pure CO2 gas with feed pressure successively increasing and then reducing, as shown in Figure 14. It was revealed that TR polymer membranes displayed enhanced property of resistance to plasticization according to the almost coincided curves of depressurization loop. The plasticization pressures of PHAA 52

and TR-PBOA membranes were listed in Table 5. The plasticization pressures of PHAA (5:5) and (10:0) were around 2.0 Mpa. After thermal treatment, the plasticization pressures were significantly increased. Especially, the sample TR 350-1(5:5) didn’t show any signs of plasticization up to 2.9 Mpa, higher than those of polyetherimide, P84 and BPZ-polycarbonate. While the plasticization of TR 400-1(10:0) occurred at lower pressure of 2.5 Mpa, which was ascribed to its larger interchain spacing despite its more rigid chain structure. The pressure dependence of the permeability of CO2 in TR polymer membranes was typical of glassy polymers. The permeability decreased with pressure and then slightly upswing. As described by dual-mode sorption and diffusion mechanism [54], the total penetrant concentration in a glassy polymer is viewed as a sum of sorption in dense region and sorption in microvoids or packing defects: ′ C = k d p + CH bp/(1 + bp),

7

′ where kd is Henry’s law solubility coefficient, CH is the Langmuir

sorption capacity, and b is the hole affinity constant. So the total solubility coefficient ( S = C/p ) decreased with pressure before plasticization occurred. At higher pressure where microvoid sorption sites became saturated, the contribution of dense regions to the overall solubility increased, resulting in significant enhancement in the effective diffusion, because penetrants in microvoids had much lower diffusivities 53

than those in normally dissolved regions. As plasticization occurred, the sorbed gases of densely packed regions resulted in a swelling of the polymer matrix with a corresponding of improvement in local chain mobility, which led to the further increase in diffusivity. The increase in diffusivity offset the decrease in solubility with pressure, thus leading to an increase in permeability above plasticization pressure. According to Figure 14 and Figure 15, above the plasticization pressure, gas permeabilities maintained almost constant and the diffusivity didn’t show a remarkable increase, indicating that TR-PBOA membranes possessed good resistance to chain motion.

(a) 70 PHAA (5:5) CO2 permeability (barrer)

60

300-1 (5:5)

50

350-1 (5:5)

40 30 20 10 0 0

0.4

0.8

1.2

1.6 2 2.4 Pressure (MPa)

2.8

3.2

3.6

(b) 54

PHAA (10:0)

CO2 permeability (barrer)

240

300-1 (10:0) 200

400-1 (10:0)

160 120 80 40 0 0

0.4

0.8

1.2

1.6 2 2.4 Pressure (MPa)

2.8

3.2

3.6

Figure 14 CO2 permeation isotherms at 35℃ for PHAA (5:5) and (10:0) at different thermal treatment temperature (blank figures represented depressurization process).

1.E-07

PHAA (5:5) CO2 diffusivity (cm2/s)

1.E-07

300-1 (5:5)

350-1 (5:5)

8.E-08 6.E-08 4.E-08 2.E-08 0.E+00 0

0.4

0.8

1.2

1.6 2 2.4 Pressure (MPa)

2.8

3.2

3.6

55

3.E-07

CO2 diffusivity (cm3 s-1)

PHAA (10:0) 3.E-07

300-1 (10:0) 400-1 (10:0)

2.E-07 2.E-07 1.E-07 5.E-08 0.E+00 0

0.4

0.8

1.2

1.6 2 2.4 Pressure (MPa)

2.8

3.2

3.6

Figure 15 CO2 diffusion coefficients for PHAA (5:5) and (10:0) at different thermal treatment temperature (blank figures represented depressurization process).

CO2 solubility (cm3(STP)cm-3 cmHg-1)

0.2 PHAA (5:5) 300-1 (5:5)

0.16

350-1 (5:5) 0.12

0.08

0.04

0 0

0.4

0.8

1.2

1.6 2 2.4 Pressure (MPa)

2.8

3.2

3.6

56

CO2 solubility (cm3(STP) cm-3 cmHg-1)

0.25

PHAA (10:0) 300-1 (10:0)

0.20

400-1 (10:0) 0.15 0.10 0.05 0.00 0

0.4

0.8

1.2

1.6 2 2.4 Pressure (MPa)

2.8

3.2

3.6

Figure 16 CO2 solubility coefficients for PHAA (5:5) and (10:0) at different thermal treatment temperature (blank figures represented depressurization process).

4 Conclusion

In this study, poly(hydroxyamide-co-amide)s (PHAAs) were derived from

commercially

available,

2,2'-bis(3-amino-4-hydroxyphenyl)

cheap

monomers:

hexafluroprop

(BisAPAF),

4,4'-diaminodiphenyl ether (ODA) and terephthaloyl chloride (TPC). Then PHAA precursor membranes underwent thermal rearrangement in the solid state at different temperature to investigate the effect of precursor structure and thermal treatment conditions on physical properties as well as thermal conversion behavior of the resulting thermally

rearranged

poly(benzoxazole-co-amide)s

(TR-PBOAs)

membranes. Thermal rearrangement process showed great dependence on glass transition temperature (Tg) of precursors, as was confirmed that 57

thermal rearrangement mainly occurred near or above glass transition temperature where chain segments obtained enough motion ability. During thermal conversion process, the Tg of TR-PBOA was increased as more rigid chains formed, preventing structure rearrangement, but the rearrangement could be restarted at higher temperature as chain segment motion obtained enough energy. Therefore, improvement in permeability was much more dependent on thermal treatment temperature rather than dwelling time. Below the Tg, thermal rearrangement occurred before chain segments received sufficient motion ability, so chain disruption in stiff chains resulting in the formation of effective micropores. Above the Tg, pore formation deriving from thermal rearrangement was partly offset by chain annealing. Furthermore, the incorporation of flexible ether linkages into backbones allowed for rearrangement at lower temperature and enhanced mechanical properties of TR-PBOA membranes. The absence of rotational ether group also increased chain packing, thus lowering FFV and gas permeability. Compared to other TR-β-PBO membranes, TR-PBOA membranes in this study exhibited superior separation performance for CO2/CH4 separation, with many samples surpassing 1991 upper bound. The ultrahigh permeability of TR-PBOA membranes was mainly ascribed to the diffusivity increment, as the coalescence of smaller cavities resulting in the higher amount of interconnected free volume with larger cavity size. 58

While, the decreased ideal selectivity derived from their reduced solubility. What’s more, these membranes displayed excellent property of resistance to CO2 plasticization according to the almost coincided curves of depressurization loop, which was essential in practical application. To realize the application of TR polymer membrane in industry, much effort should be focused on the enhancement of gas transport performance and the preparation of hollow fiber membranes in the future.

Acknowledgments

Financial support is gratefully acknowledged from the National Natural Science Foundation of China (No. 21436009)

References [1] D.E. Sanders, Z.P. Smith, R.L. Guo, L.M. Robeson, J.E. McGrath, D.R. Paul, B.D. Freeman, Energy-efficient polymeric gas separation membranes for a sustainable future: A review, Polymer, 54 (2013) 4729-4761. [2] I. Sreedhar, R. Vaidhiswaran, B.M. Kamani, A. Venugopal, Process and engineering trends in membrane based carbon capture, Renewable and Sustainable Energy Reviews, 68 (2017) 659-684. [3] X.Q. Cheng, Z.X. Wang, X. Jiang, T. Li, C.H. Lau, Z. Guo, J. Ma, L. Shao, Towards sustainable ultrafast molecular-separation membranes: From conventional polymers to emerging materials, Progress in Materials Science, 92 (2018) 258-283. [4] I. Pinnau, L.G. Toy, Transport of organic vapors through poly(1-trimethylsilyl-1-propyne), Journal of Membrane Science, 116 (1996) 199-209. [5] L.M. Robeson, M.E. Dose, B.D. Freeman, D.R. Paul, Analysis of the transport properties of thermally rearranged (TR) polymers and polymers of intrinsic microporosity (PIM) relative to upper bound performance, Journal of Membrane Science, 525 (2017) 18-24. [6] L.M. Robeson, The upper bound revisited, Journal of Membrane Science, 320 (2008) 390-400. [7] Z.X. Low, P.M. Budd, N.B. McKeown, D.A. Patterson, Gas Permeation Properties, Physical Aging, and Its Mitigation in High Free Volume Glassy Polymers, Chemical reviews, 118 (2018) 5871-5911. [8] P.M. Budd, N.B. McKeown, B.S. Ghanem, K.J. Msayib, D. Fritsch, L. Starannikova, N. Belov, O. 59

Sanfirova, Y. Yampolskii, V. Shantarovich, Gas permeation parameters and other physicochemical properties of a polymer of intrinsic microporosity: Polybenzodioxane PIM-1, Journal of Membrane Science, 325 (2008) 851-860. [9] R. Swaidan, M. Al-Saeedi, B. Ghanem, E. Litwiller, I. Pinnau, Rational Design of Intrinsically Ultramicroporous Polyimides Containing Bridgehead-Substituted Triptycene for Highly Selective and Permeable Gas Separation Membranes, Macromolecules, 47 (2014) 5104-5114. [10] N.B. McKeown, P.M. Budd, Polymers of intrinsic microporosity (PIMs): organic materials for membrane separations, heterogeneous catalysis and hydrogen storage, Chem Soc Rev, 35 (2006) 675-683. [11] Y. Yin, M.D. Guiver, Microporous polymers: Ultrapermeable membranes, Nat Mater, 16 (2017) 880-881. [12] G. Dong, Y.M. Lee, Microporous polymeric membranes inspired by adsorbent for gas separation, Journal of Materials Chemistry A, 5 (2017) 13294-13319. [13] D.F. Sanders, Z.P. Smith, C.P. Ribeiro, R. Guo, J.E. McGrath, D.R. Paul, B.D. Freeman, Gas permeability, diffusivity, and free volume of thermally rearranged polymers based on 3,3 ′ -dihydroxy-4,4 ′ -diamino-biphenyl (HAB) and 2,2 ′ -bis-(3,4-dicarboxyphenyl) hexafluoropropane dianhydride (6FDA), Journal of Membrane Science, 409-410 (2012) 232-241. [14] S. Kim, Y.M. Lee, Rigid and microporous polymers for gas separation membranes, Progress in Polymer Science, 43 (2015) 1-32. [15] H.B. Park, C.H. Jung, Y.M. Lee, A.J. Hill, S.J. Pas, S.T. Mudie, E. Van Wagner, B.D. Freeman, D.J. Cookson, Polymers with Cavities Tuned for Fast Selective Transport of Small Molecules and Ions, Science, 318 (2007) 254-258. [16] C.H. Park, E. Tocci, Y.M. Lee, E. Drioli, Thermal treatment effect on the structure and property change between hydroxy-containing polyimides (HPIs) and thermally rearranged polybenzoxazole (TR-PBO), J Phys Chem B, 116 (2012) 12864-12877. [17] S.H. Han, N. Misdan, S. Kim, C.M. Doherty, A.J. Hill, Y.M. Lee, Thermally Rearranged (TR) Polybenzoxazole: Effects of Diverse Imidization Routes on Physical Properties and Gas Transport Behaviors, Macromolecules, 43 (2010) 7657-7667. [18] Y.S. Do, W.H. Lee, J.G. Seong, J.S. Kim, H.H. Wang, C.M. Doherty, A.J. Hill, Y.M. Lee, Thermally rearranged (TR) bismaleimide-based network polymers for gas separation membranes, Chemical communications, 52 (2016) 13556-13559. [19] M. Calle, Y.M. Lee, Thermally Rearranged (TR) Poly(ether−benzoxazole) Membranes for Gas Separation, Macromolecules, 44 (2011) 1156-1165. [20] M. Calle, Y. Chan, H.J. Jo, Y.M. Lee, The relationship between the chemical structure and thermal conversion temperatures of thermally rearranged (TR) polymers, Polymer, 53 (2012) 2783-2791. [21] R. Guo, D.F. Sanders, Z.P. Smith, B.D. Freeman, D.R. Paul, J.E. McGrath, Synthesis and characterization of Thermally Rearranged (TR) polymers: influence of ortho-positioned functional groups of polyimide precursors on TR process and gas transport properties, J. Mater. Chem. A, 1 (2013) 262-272. [22]

M.

Calle,

C.M.

Doherty,

A.J.

Hill,

Y.M.

Lee,

Cross-Linked

Thermally

Rearranged

Poly(benzoxazole-co-imide) Membranes for Gas Separation, Macromolecules, 46 (2013) 8179-8189. [23] M. Calle, H.J. Jo, C.M. Doherty, A.J. Hill, Y.M. Lee, Cross-Linked Thermally Rearranged Poly(benzoxazole-co-imide) Membranes Prepared from ortho-Hydroxycopolyimides Containing Pendant Carboxyl Groups and Gas Separation Properties, Macromolecules, 48 (2015) 2603-2613. 60

[24] S. Li, H.J. Jo, S.H. Han, C.H. Park, S. Kim, P.M. Budd, Y.M. Lee, Mechanically robust thermally rearranged (TR) polymer membranes with spirobisindane for gas separation, Journal of Membrane Science, 434 (2013) 137-147. [25] Y. Zhuang, J.G. Seong, W.H. Lee, Y.S. Do, M.J. Lee, G. Wang, M.D. Guiver, Y.M. Lee, Mechanically Tough, Thermally Rearranged (TR) Random/Block Poly(benzoxazole-co-imide) Gas Separation Membranes, Macromolecules, 48 (2015) 5286-5299. [26] Q. Liu, D.R. Paul, B.D. Freeman, Gas permeation and mechanical properties of thermally rearranged (TR) copolyimides, Polymer, 82 (2016) 378-391. [27] S. Kim, J. Hou, Y. Wang, R. Ou, G.P. Simon, J.G. Seong, Y.M. Lee, H. Wang, Highly permeable thermally rearranged polymer composite membranes with a graphene oxide scaffold for gas separation, Journal of Materials Chemistry A, 6 (2018) 7668-7674. [28] K.T. Woo, G. Dong, J. Lee, J.S. Kim, Y.S. Do, W.H. Lee, H.S. Lee, Y.M. Lee, Ternary mixed-gas separation for flue gas CO 2 capture using high performance thermally rearranged (TR) hollow fiber membranes, Journal of Membrane Science, 510 (2016) 472-480. [29] H. Wang, D.R. Paul, T.-S. Chung, The effect of purge environment on thermal rearrangement of ortho-functional polyamide and polyimide, Polymer, 54 (2013) 2324-2334. [30] A. Kushwaha, M.E. Dose, Z.P. Smith, S. Luo, B.D. Freeman, R. Guo, Preparation and properties of polybenzoxazole-based gas separation membranes: A comparative study between thermal rearrangement (TR) of poly(hydroxyimide) and thermal cyclodehydration of poly(hydroxyamide), Polymer, 78 (2015) 81-93. [31] H. Borjigin, Q. Liu, W. Zhang, K. Gaines, J.S. Riffle, D.R. Paul, B.D. Freeman, J.E. McGrath, Synthesis and characterization of thermally rearranged (TR) polybenzoxazoles: Influence of isomeric structure on gas transport properties, Polymer, 75 (2015) 199-210. [32] H.J. Jo, C.Y. Soo, G. Dong, Y.S. Do, H.H. Wang, M.J. Lee, J.R. Quay, M.K. Murphy, Y.M. Lee, Thermally Rearranged Poly(benzoxazole-co-imide) Membranes with Superior Mechanical Strength for Gas Separation Obtained by Tuning Chain Rigidity, Macromolecules, 48 (2015) 2194-2202. [33] C.Y. Soo, H.J. Jo, Y.M. Lee, J.R. Quay, M.K. Murphy, Effect of the chemical structure of various diamines on the gas separation of thermally rearranged poly(benzoxazole-co-imide) (TR-PBO-co-I) membranes, Journal of Membrane Science, 444 (2013) 365-377. [34] C. Duan, X. Jie, D. Liu, Y. Cao, Q. Yuan, Post-treatment effect on gas separation property of mixed matrix membranes containing metal organic frameworks, Journal of Membrane Science, 466 (2014) 92-102. [35]

Y.S.

Do,

J.G.

Seong,

S.

Kim,

J.G.

Lee,

Y.M.

Lee,

Thermally

rearranged

(TR)

poly(benzoxazole-co-amide) membranes for hydrogen separation derived from 3,3′-dihydroxy-4,4′ -diamino-biphenyl (HAB), 4,4 ′ -oxydianiline (ODA) and isophthaloyl chloride (IPCl), Journal of Membrane Science, 446 (2013) 294-302. [36] B. Comesaña-Gándara, A. Hernández, J.G. de la Campa, J. de Abajo, A.E. Lozano, Y.M. Lee, Thermally rearranged polybenzoxazoles and poly(benzoxazole-co-imide)s from ortho-hydroxyamine monomers for high performance gas separation membranes, Journal of Membrane Science, 493 (2015) 329-339. [37] G. Dong, K.T. Woo, J. Kim, J.S. Kim, Y.M. Lee, Simulation and feasibility study of using thermally rearranged polymeric hollow fiber membranes for various industrial gas separation applications, Journal of Membrane Science, 496 (2015) 229-241. [38] H. Wang, T.-S. Chung, The evolution of physicochemical and gas transport properties of thermally 61

rearranged polyhydroxyamide (PHA), Journal of Membrane Science, 385-386 (2011) 86-95. [39] A.C. Comer, C.P. Ribeiro, B.D. Freeman, S. Kalakkunnath, D.S. Kalika, Dynamic relaxation characteristics of thermally rearranged aromatic polyimides, Polymer, 54 (2013) 891-900. [40] S.H. Han, H.J. Kwon, K.Y. Kim, J.G. Seong, C.H. Park, S. Kim, C.M. Doherty, A.W. Thornton, A.J. Hill, A.E. Lozano, K.A. Berchtold, Y.M. Lee, Tuning microcavities in thermally rearranged polymer membranes for CO2 capture, Phys Chem Chem Phys, 14 (2012) 4365-4373. [41] B. Comesaña-Gándara, M. Calle, H.J. Jo, A. Hernández, J.G. de la Campa, J. de Abajo, A.E. Lozano, Y.M. Lee, Thermally rearranged polybenzoxazoles membranes with biphenyl moieties: Monomer isomeric effect, Journal of Membrane Science, 450 (2014) 369-379. [42] . [43] L.M. Robeson, CORRELATION OF SEPARATION FACTOR VERSUS PERMEABILITY FOR POLYMERIC MEMBRANES, Journal of Membrane Science, 62 (1991) 165-185. [44] H. Wang, S. Liu, T.-S. Chung, H. Chen, Y.-C. Jean, K.P. Pramoda, The evolution of poly(hydroxyamide amic acid) to poly(benzoxazole) via stepwise thermal cyclization: Structural changes and gas transport properties, Polymer, 52 (2011) 5127-5138. [45] S. Luo, J. Liu, H. Lin, B.A. Kazanowska, M.D. Hunckler, R.K. Roeder, R. Guo, Preparation and gas transport properties of triptycene-containing polybenzoxazole (PBO)-based polymers derived from thermal rearrangement (TR) and thermal cyclodehydration (TC) processes, Journal of Materials Chemistry A, 4 (2016) 17050-17062. [46] S. Neyertz, D. Brown, S. Pandiyan, N.F.A. van der Vegt, Carbon Dioxide Diffusion and Plasticization in Fluorinated Polyimides, Macromolecules, 43 (2010) 7813-7827. [47] T.C. Merkel, V.I. Bondar, K. Nagai, B.D. Freeman, I. Pinnau, Gas sorption, diffusion, and permeation in poly(dimethylsiloxane), Journal of Polymer Science Part B-Polymer Physics, 38 (2000) 415-434. [48] [49] T.C. Merkel, V. Bondar, K. Nagai, B.D. Freeman, Sorption and transport of hydrocarbon and perfluorocarbon gases in poly(1-trimethylsilyl-1-propyne), Journal of Polymer Science Part B-Polymer Physics, 38 (2000) 273-296. [50] S. Kim, J.G. Seong, Y.S. Do, Y.M. Lee, Gas sorption and transport in thermally rearranged polybenzoxazole membranes derived from polyhydroxylamides, Journal of Membrane Science, 474 (2015) 122-131. [51] W.L. A.F. Ismail, Penetrant-induced plasticization phenomenon in glassy polymers for gas separation membrane, (2002). [52] M.R. Coleman, W.J. Koros, Conditioning of fluorine-containing polyimides. 2. Effect of conditioning protocol at 8% volume dilation on gas-transport properties, Macromolecules, 32 (1999) 3106-3113. [53] A. Bos, I.G.M. Punt, M. Wessling, H. Strathmann, Suppression of CO2-plasticization by semiinterpenetrating polymer network formation, Journal of Polymer Science Part B-Polymer Physics, 36 (1998) 1547-1556. [54] T.A. Barbari, W.J. Koros, D.R. Paul, GAS-TRANSPORT IN POLYMERS BASED ON BISPHENOL-A, Journal of Polymer Science Part B-Polymer Physics, 26 (1988) 709-727.

62

Highlights  TR-PBOA copolymer membranes were prepared by in-situ thermal treating PHAA precursors.  Thermal rearrangement mainly occurred near or above glass transition temperature.  The effect of dwell time on thermal conversion ratio was minor.  TR-PBOA membranes displayed superior CO2/CH4 separation performance.

63