Climate change, induced by steadily increasing CO₂ emission since the 1900s, has emerged as a global challenge. There is an urgent need to adopt active measures to mitigate carbon emission[1–3]. The European Commission has identified carbon capture and storage as a key decarbonization option for the electricity sector and energy-intensive industries and has also pledged the achievement of carbon neutrality by 2050. Furthermore, the net worth of global carbon capture and storage market is expected to reach 6.13 × 10⁹ USD per annum by 2027[4–6]. Therefore, developing superior separation technologies to capture or purify exhaust gases is necessary to sustain national prosperity. In this regard, membrane technology has emerged as a promising alternative to CO₂ capture because of its high energy efficiency, low footprint, and modifiable operation process[7–9]. Particularly, CO₂-philic membranes, fabricated with poly(ethylene oxide) (PEO) and its derivatives, have been widely studied for their strong dipole– quadrupole interactions with CO₂ molecules[10,11]. However, the trade-off between permeability and selectivity significantly hinders the separation performance of traditional PEO-containing polymers and limits their extensive application in industrial processes [12]. Consequently, it is necessary to further explore the potential of PEO-based materials.

The conventional strategy to address this trade-off is via the preparation of mixed matrix membranes (MMMs), which consist of two parts: a processable polymer matrix and highly efficient inorganic particles (as fillers)[13,14]. However, the combination of organic and inorganic materials easily leads to the uneven dispersion of fillers and undesirable defects between polymer matrix and inorganic materials[15–17]. Considering the uncertain compatibility between polymers and fillers, to overcome the limitations of MMMs, polymer blending has been adopted owing to its convenience, implementation ability, and commercial potential[18–21]. Low molecular weight PEOs (LMWPEO) are a class of highly efficient additives, which not only allows mutual miscibility with PEO-based materials but also improves polymer flexibility, optimizes fractional free volumes (FFVs), and eventually promotes gas transport[22,23]. Yave et al. [24] reported that when polyethylene glycol dimethyl ether (PEGDME) was employed as an additive to the Pebax matrix, the phase separation of PEO and polyamide 6 (PA6) was enhanced by the alkyl groups on PEGDME. This resulted in a larger number of PEO chain segments forming microdomains that were more conducive to CO₂ transport. Moreover, the glass transition temperature (T_g) of the polymer was reduced from 220 to 195 K, providing further evidence of improved chain segment motility. Therefore, the CO₂ permeability of the Pebax-PEGDME membrane is much higher than that of the pristine Pebax membrane. Furthermore, our group was the first to demonstrate that the gas transport performance of cross-linked PEO membranes can be improved via a facile impregnation method [25]. FFVs of the cross-linked PEO films consistently increased with the PEG amount, owing to the plasticization effect. Additionally, the immersed PEG significantly enhanced the CO₂ solubility of the membrane as a function of its high ethylene oxide (EO) content. 

However, embedded LMWPEOs cause structural instability. The mechanical strength of a membrane primarily relies on the entanglement and crystal of polymeric chains[26,27]. Short linear PEO molecules weaken the interactions among adjacent chains. Therefore, the amount of additives that can be used is significantly limited because of the deterioration of mechanical property of the membrane, hindering further improvement of its separation performance. What’s more, LMWPEOs are also easily extruded at high operating pressures, damaging the membrane structure and causing separation performance degradation, which limits their application under harsh conditions. Hence, addressing this trade-off between mechanical property and separation efficiency is of critical urgency

Herein, we propose a one-step method that incorporates poly(ethylene glycol) methyl ether acrylate (PEGMEA) with high EO content into the Pebax matrix. The in situ polymerization of vinyl-functionalized PEGMEA monomers can be initiated via thermal treatment, while the methyl termination of the side chains can create an open structure for the ease of transport. The resulting highly branched poly-PEGMEA helps generate sufficient entanglement with polymer chains to guarantee adequate mechanical strength even under massive loading. This self-polymerized membrane is expected to subtly synergize the robust physical structure with the gas separation performance, unlike a simply blended Pebax/PEGMEA membrane. We believe this study will open a pathway for the rational design of ultra-stable membranes under harsh conditions.

Pebax^® MH 1657 (Pebax) was obtained from Arkema Co., Ltd. (France). PEGMEA (number-average molecular weight (M_n) = 480) and 2,2' -azobis (2-methylpropionitrile) (AIBN) were obtained from Sigma-Aldrich Co., LLC (USA). Anhydrous ethanol (AR) was supplied by Tianjin Kemao Chemical Reagent Co., Ltd. (China). Deionized water (DI) was used throughout the study. 

A specified quantity of Pebax was dissolved in a mixture of 70 wt% ethanol and 30 wt% water. Thereafter, the obtained mixture was refluxed for 2 h at 80 °C to acquire 3 wt% Pebax homogeneous solution. After cooling to ambient temperature, the solution was cast on a homemade glass mold, dried in air for 48 h, and then heated in a vacuum oven at 40 °C to ensure complete solvent evaporation. A self-polymerization-confined membrane was fabricated by uniformly mixing PEGMEA in a Pebax solution, while AIBN (1 wt%) was simultaneously added as an initiator. After solvent evaporation, the glass mold was heated to 80 °C in a vacuum oven for 2 h to ensure sufficient polymerization. For the ease of differentiation, the resulting self-polymerized membranes were denoted as SPM-A, and the membrane without further thermal treatment was denoted as PMEA-B, where A and B are the mass fractions of PEGMEA. The mass fraction (w_f) was calculated using Eq. (1):

where m₁ and m₂ represent the masses of PEGMEA and Pebax matrix, respectively.

Attenuated total reflectance Fourier transform infrared (ATRFTIR) spectra were obtained using a PerkinElmer Spectrum Two (USA) at ambient temperature from 4000 to 500 cm^–1 . Packing motifs of polymer chains were measured using wide-angle X-ray diffraction (WAXD) on a Bruker D8 ADVANCE X-ray diffractometer (40 kV, 40 mA, Cu Kα, = 1.5418 Å; Germany) in a 5°–80° range with a scan speed of 5°·min^–1 at room temperature. X-ray photoelectron spectroscopy (XPS) was conducted on a Thermo ESCALAB 250XI (Al Kα, = 1486.6 eV; Thermo Scientific, USA) at room temperature. Differential scanning calorimetry (DSC; Q200, TA) was used to study thermal properties. Initial heating was performed at a heating rate of 20 °C·min^–1 to eliminate the heating history. The exothermic curves of the membranes were recorded from –90 to 250 °C at a scanning rate of 10 °C·min^–1 under a N₂ atmosphere. The area swelling and water uptake of the membranes were determined by soaking the dry film (cut into 2 cm diameter round pieces) in water for 72 h. Water was changed every 12 h. The ratio of the area difference between the dry and wet films to that of the dry membrane corresponds to the area swelling of the films. Water uptake was calculated in the same manner. Weight loss was measured as the ratio of the weight difference between the membranes before and after water absorption to the dry membrane before water was absorbed. The membranes were trimmed into 4 cm × 1 cm pieces for tensile testing (tensile testing machine was CTM2050; Xie Qiang, China), and the stretching velocity was set at 5 mm·min^–1 . The membrane thickness was measured using a micrometer caliper, and each membrane was measured nine times to obtain an average value. Small-angle X-ray scattering (SAXS) was performed via Saxesess mc2 (Anton Paar, Germany) using Kα radiation from 0.08° to 5° at room temperature. The thermal stability of membranes was tested using a thermogravimetric analyzer (TGA; Q500, TA) at a heating rate of 10 °C·min^–1 from 20 to 800 °C under a N₂ atmosphere. Polymer density was determined via the buoyancy method using a Mettler Toledo balance (ME104; Switzerland) and density determination kit. The membrane density (ρ_m) was calculated as follows: 

where m_a and m_b represent membrane weights in air and auxiliary liquids, respectively. ρ₀ is the density of the silicone oil auxiliary liquid (0.968 g·cm^–3 ).

as permeation was analyzed using a homemade apparatus employing the constant volume/variable pressure method. Following our previous study, gas permeation was tested in the following order: H₂ → N₂ → CO₂ [28]. The gas pressure was maintained at 3.5 bar (1 bar = 10⁵ Pa), and the test temperature was controlled at 35 °C, unless specified otherwise. Detailed information on gas testing is provided in Appendix A.

SPM membranes were obtained by introducing PEGMEA in the Pebax solution, followed by in situ polymerization during thermal treatment (Fig. 1). Poly-PEGMEA exhibits a highly branched structure, which is capable of dynamic entanglement with polymer chains. For the PMEA membrane, the chain length of the low molar weight PEGMEA is very short to form an effective entanglement. Therefore, the maximum PEGMEA loading of the PMEA membrane was 70 wt%, while it was up to 90 wt% for SPM membranes, which represents the state-of-the-art in similar membrane materials. 

Fig. 1. Schematic illustration of the synthesis process of SPM membranes.

To confirm the construction and reliability of the selfpolymerized structure, we measured the weight loss, water uptake, and area swelling of the SPM and PMEA membranes by soaking them in pure water for 72 h (Figs. 2(a) and (b); Figs. S1 and S2 in Appendix A). As shown in Fig. 2(a), almost all SPM membranes exhibited a gel content of more than 80% in water, demonstrating that the linear PEGMEA was successfully polymerized and the resulting poly-PEGMEA intertwined strongly with Pebax chains. In contrast, the weight loss of the PMEA membranes increased with PEGMEA loading, which proved that most of the PEGMEA content dissolved out from the membranes. The opposite trend in water absorption reflected the loss of the EO content in the PMEA membrane, whereas the SPM-90 membrane showed approximately three times larger absorption capability than the Pebax membrane. This is because of the introduction of a large number of hydrophilic EO groups. Similarly, the distinct difference in the area swelling (Figs. S1 and S2) also indicated the formation of a new stable structure that endows the membrane with high stability.

Mechanical properties of the membrane are represented by the stress–strain curves shown in Figs. 2(c) and (d) and Fig. S3 in Appendix A. The incremental embedding of PEGMEA deteriorates the mechanical strength of the Pebax matrix owing to the plasticization effect. On the contrary, the SPM membranes exhibit a higher performance than the PMEA membranes in terms of both breaking strength and maximum elongation, confirming that the thermal treatment successfully intertwined Pebax chains and branched poly-PEGMEA. This dynamic entanglement facilitates a firm interlocking in the self-polymerized structure to reinforce the polymer chains and prevent the membrane from breaking. In other words, self-polymerization-confined membranes show better mechanical stability than conventional LMWPEO composite materials, making the membranes suitable for a wider range of industrial applications.

Fig. 2. Comparison of macroscopic properties between PMEA and SPM membranes. (a) Weight loss and (b) water uptake of membranes after water soaking for 72 h. (c, d) Stress–strain curves.

Chemical structures of the SPM membranes were analyzed using XPS and ATR-FTIR. According to the XPS results (Fig. 3(a); Fig. S4 and Table S1 in Appendix A), the disappearance of the C–N peak at 286.05 eV and the diminution of N atom percentage confirm the introduction of PEGMEA. The peak at 283.3 eV in the C 1s spectrum of SPM-90 corresponds to aliphatic chains, confirming the formation of poly-PEGMEA[29–31]. In the ATR-FTIR spectra of Pebax (Fig. S5 in Appendix A), the absorption intensity of the amide N–H (3299 cm^–1 ) and amide carbonyl H–N–C=O (1640 cm^–1 ) groups decreased. On the contrary, the absorption intensity of C=O groups (1730 cm^–1 ) in both Pebax and PEGMEA increased with the addition of PEGMEA[27,32,33]. The strongest adsorption bond at approximately 1100 cm^–1 denotes the asymmetrical stretching vibration of C–O–C units [34]. No significant band shift was observed between the different Fourier transform infrared (FT-IR) spectra, indicating a similar chemical structure and good compatibility of these two polymers, which is beneficial for the construction of a homogeneous membrane.

The FFV of a polymer is related to its density. Herein, we measured the membrane density with different PEGMEA volume fractions and compared them with the results of the additive model, as shown in Fig. 3(b). The additive model is defined by the following Eq. (3):

where Φ₁ is the Pebax volume fraction, which can be estimated using the experimental density. ρ₁ is the Pebax density, and ρ₂ represents the density of PEGMEA and poly-PEGMEA in the PMEA and SPM membranes, respectively. The density of poly-PEGMEA was adopted from the Freeman’s study [35]. Notably, the density of the composite membrane merely relies on the polymer component of the system, according to Eq. (3). Therefore, the distinction in density between the experimental data and theoretical values may exist owing to the structural variation of the polymer (i.e., variation of FFV).

As shown in Fig. 3(b), the values for SPM membranes are higher than those for PMEA membranes, indicating the selfpolymerization of PEGMEA in SPM membranes, which is in agreement with XPS results. In addition, the experimental densities of both SPM and PMEA membranes were lower than those obtained using the additive model. Specifically, when the volume fraction of PEGMEA is less than 50%, the density of the SPM membrane is even less than that of the PMEA membrane obtained using the additive model. This proves that poly-PEGMEA can effectively disturb the polymer structure of Pebax. Pebax is a typical copolymer consisting of a PEO and PA6 segments, and a distinct hydrogen bond interaction exists between the amide bond and the EO unit. In SPM membranes, poly-PEGMEA might be inserted between the PEO and PA6 chains, and the methoxy end group would prohibit any hydrogen bonding interactions to prevent the assembly of PEO and PA6 chains. Therefore, the resulting hybrid polymer has a large intermolecular space for micro-Brownian motion. This speculation is also supported by the decrease in the glass transition temperature, which will be discussed in the following section.

Fig. 3. Structural characterization of membranes. (a) C 1s spectra. (b) Density. (c) DSC thermograms. (d) SAXS patterns, where the q represents for the scattering vector.

Thermal properties of SPM membranes were investigated using DSC (Fig. 3(c)). There are two distinct endothermic peaks at 16.5 and 205.5 °C, which correspond to the melting points of the PEO and PA6 segments, respectively, indicating microphase separation in the Pebax membrane. The degree of crystallinity (X_C) can be obtained using the melting enthalpy (Table 1) [20]. The X_C of PEO increases with the increase in PEGMEA loading (higher EO amounts). On the contrary, for the PA6 phase, the low molecular weight PEGMEA deteriorates the crystal structure with small and imperfect crystallites, which lead to a decreased melting temperature (T_m) and X_C. Although the crystallinity of the PEO phase gradually increases with the addition of PEGMEA, T_m decreases from 16.5 to 4.7 °C, indicating that the PEO phase is in the molten state under conventional operating temperature conditions. Therefore, the amorphous proportion in the blend polymer increases. Besides, the T_g of SPM membranes decreases progressively from –49.1 to –58.5 °C with increased loading. This is because PEGMEA acts as a plasticizer for the Pebax matrix, which can reinforce chain mobility, as discussed above. The change in T_g can be associated with the FFV of the polymer [36]. The FFV increases with a reduction in the T_g for rubbery polymers. Therefore, it is possible that gas diffusivity and solubility improve as the amorphous proportion and free volume increases. Furthermore, all SPM membranes remained stable up to 300 °C, according to the TGA results (Fig. S6 in Appendix A), suggesting their suitability for high-temperature applications.

Table 1 Thermal properties of the Pebax and SPM membranes.

SAXS is an effective method to investigate the microstructure of polymers and can directly reflect the microphase separation of block copolymers. All SAXS curves (Fig. 3(d)) exhibit a broad peak, indicating that the membranes form microdomains with a shortrange order. Pristine Pebax exhibits a micro-phase separation structure with a nanodomain size of 16.4 nm, according to the Bragg relation [37]. As the PEGMEA content increased, a gradual rightward shift of the peak position was observed, and the intensity gradually decreased, suggesting a breakdown of the microphase separation structure. Similarly, according to the WAXD curves (Fig. S7 in Appendix A), the intensity of the PA6 peak at 2θ = 24° gradually decreased. This suggests that poly-PEGMEA adversely affects the crystallization behavior of the PA6 phase, leading to reduced crystallinity and an increase in the amorphous domain in the SPM membrane, thereby corroborating the DSC and SAXS results. 

Figs. 4(a) and (b) show the pure gas separation performances of the Pebax and SPM membranes at 3.5 bar and 35 °C. As the PEGMEA content increased from 0 to 90%, the permeability of all gases improved significantly. Specifically, the CO₂ permeability increased considerably by 568% from 100.6 Barrer (1 Barrer = 10^–10 cm³ (STP)·cm·(cm² ·s·cmHg)^–1 = 7.5006 × 10^–18 m³ ·m·(m² ·s·Pa)^–1 ) for pristine Pebax to 672.4 Barrer for the SPM-90 membrane. Considering this together with the DSC data, the increase in gas permeability is largely due to the optimization of the free volume. First, the introduction of PEGMEA alters the stacking of the original Pebax chains and enhances their motility. Second, the free unreacted –OCH₃ terminal group acts as a spacer to further expand the free volume fraction of the membrane. Moreover, PEGMEA has a higher EO content than Pebax, which allows an intensive CO₂ solution process. The CO₂/N₂ selectivity marginally increased from 44 to 47, and the CO₂/H₂ selectivity improved by 51.8% (to 12.9) with the PEGMEA incorporation. 

To clarify the reason behind the change in the separation performance after the PEGMEA embedment, gas solubility and diffusivity were studied using the time-lag method. Normalized coefficients were employed to study the effect of the PEGMEA content on the gas transport properties (Figs. 4(c) and (d); Table S2 in Appendix A). The solubility and diffusivity of all gases showed an increasing trend. Crystalline polymers usually exhibit low gas solubility owing to their regular and tightly packed structures [38]. However, the introduction of PEGMEA reduces the proportion of crystalline phases in the membrane and increases the amorphous phase, thereby improving the gas solubility of the membrane. Moreover, owing to the dipolar–quadrupole interaction, the increase in EO units in the membrane improves the specific solubility of the membrane for CO₂, ultimately enabling an overall increase in the solubility selectivity. Gas diffusivity is often related to the polymer FFV. Therefore, gas diffusivities increase with increasing PEGMEA content. Notably, when the FFV of the membrane increases, gas molecules with larger kinetic diameters are more influential, and therefore, the gas diffusion coefficient increases in the sequence of N₂ > CO₂ > H₂. This results in an increase and a decrease in CO₂/H₂ and CO₂/N₂ diffusivity selectivities, respectively. Consequently, CO₂/H₂ selectivity significantly increases, which suggests that this novel membrane is suitable for CO₂/light gas separation.

Fig. 4. (a) Gas permeability (1 Barrer = 7.5006 × 10^–18 m³ ·m·(m² ·s·Pa)^–1 ), (b) selectivity, (c) normalized solubility, and (d) normalized diffusivity of Pebax and SPM membranes at 3.5 bar and 35 °C.

Simultaneously, the Maxwell model was used to estimate gas permeation in SPM membranes, which has been proven to be valid over the entire composition range for the dispersion of isometric particles. The permeability of the composite membrane can be described using the following equation:

where P_b, P_c, and P_d are gas permeabilities of the composite material, continuous phase, and discontinuous phase, respectively. The volume fraction of the discontinuous phase is represented by Φ. Herein, P_c and P_d are the intrinsic permeabilities of Pebax and poly-PEGMEA, respectively. The value for Pd was obtained from the Freeman’s study [35]. The experimental and theoretical values of gas permeability in SPM membranes are shown in Fig. 5, and a distinct deviation between the experimental and Maxwell model values is observed, which may be because of the alteration of the structure of the Pebax matrix by the introduction of PEGMEA. Although this model provides only a crude estimation of the connection between the permeability and composition of polymer blends, the results support the conclusion that the significant enhancement of gas permeability is related to the membrane FFV changes owing to the addition of PEGMEA.

Fig. 5. Experimental values and Maxwell model estimation of the variation of CO₂ permeability with different PEGMEA volume fractions.

A pressure test was conducted to further investigate the operational stability of the membranes. As shown in Fig. 6(a), both CO₂ permeability and CO₂/N₂ selectivity of SPM-90 membrane gradually increase as the feed pressure increases from 3.5 to 20 bar. The increased CO₂ permeability is mainly due to the CO₂ solubility because high pressure can facilitate gas dissolution in a polymer. Therefore, a large amount of CO₂ in the membrane promotes the motility of polymer chains, thereby increasing the FFV, which is known as the plasticization effect. In contrast, N₂ permeability decreases monotonically (Fig. S8 in Appendix A) because the N₂ solubility is low. Plasticization effect does not occur, and membrane compression at high feed pressures leads to FFV loss [28]. Consequently, these two separate trends jointly promoted the enhancement of the CO₂/N₂ selectivity. For the PMEA-70 membrane, CO₂ permeability decreased from 916.1 to 710.3 Barrer as the pressure increased. The reason for this can be inferred from Fig. S9 in Appendix A. The linear PEGMEA was forced out of the membrane after the pressure test, which has also been reported by Shin et al. and termed as the ‘‘leaching out’’ effect [39]. Consequently, when gas pressure decreased from 20 to 3.5 bar, the separation performance of the SPM-90 membrane is reversible (644.7 Barrer), while the PMEA-70 membrane showed a significant decrease in performance to 449.3 Barrer owing to the PEGMEA loss. The enhanced performance of the SPM membrane in the pressure test emphasizes the importance of the self-polymerized structure, which allows a high PEGMEA loading while maintaining superior structural stability.

Fig. 6. Effect of (a) feed pressure and (b) temperature on the separation performance. (c) Long-term stability of SPM-90 membrane at 3.5 bar and 35 °C for 120 h. (d, e) Comparison of the separation performance of SPM membranes with an upper-bound line.

The effects of the operating temperatures on the CO₂ permeability and CO₂/N₂ selectivity are shown in Fig. 6(b) and Fig. S10 in Appendix A. The CO₂ permeability is notably improved as the temperature increases, owing to the enhancement in the average kinetic energy and mobility of the polymer chains [40]. However, a decreasing trend was observed for CO₂/N₂ selectivity, which corresponds to the decreased CO₂ solubility, as previously reported [41]. In addition, the activation energy is often used to interpret the temperature dependence of the membrane separation performance, which can be described using the Arrhenius equation: 

where P_A is the permeability (Barrer); P_A0 represents the preexponential factor; R is the universal gas constant (J·mol^–1 ·K^–1 ); T is the absolute temperature (K), and E_P is the activation energy (kJ·mol^–1 ). The calculated results are presented in Table S3 in Appendix A. Both the CO₂ and N₂ apparent activation energy values (from 30 to 55 °C) of SPM-90 are smaller than those of pristine Pebax membranes, revealing the lower gas permeation energy barrier for the SPM membrane. 

The long-term stability of the SPM-90 membrane for CO₂/N₂ separation was investigated at 3.5 bar and 35 °C for 120 h. As shown in Fig. 6(c), the CO₂ permeability slightly decreased in the first 24 h, which may be due to the loss of free PEGMEA. However, the value was finally stabilized at 650 Barrer after 120 h. Therefore, this minor change does not have a significant effect on the membrane selectivity, demonstrating the outstanding long-term stability of our SPM membranes.

Figs. 6(d) and (e) compare the CO₂/N₂ and CO₂/H₂ separation performances of the SPM membranes with Robeson’s upperbound line [47]. As the PEGMEA content increased, the separation performance of the SPM membranes approached the upper limit, while the performance of the SPM-90 membrane surpassed the upper-bound limit for the CO₂/N₂ gas pair. As for CO₂/H₂ separation, our study demonstrated a simultaneous increase in both permeability and selectivity, thereby mitigating the well-known ‘‘trade-off” effect. Moreover, a higher test pressure increased the membrane performance to a higher level, which is beneficial for the practical application of CO₂-philic membranes.

Finally, the comprehensive performance of our SPM membranes was competitive with that of other Pebax-based polymer blending membranes (Table 2) [19,20,24,32,39,42-46], outlining the nature of the self-polymerized structure. The SPM additive content is higher than the previously reported values and can reach as high as 90 wt%, benefitting from sufficient entanglement between poly-PEGMEA and Pebax chains. Meanwhile, the large number of EO units increases the affinity of the membrane for CO₂ molecules and enables a superior separation performance compared with other polymer blending membranes especially driven by the plasticizing effect. Therefore, we believe that our self-polymerization-confined all-polymeric CO₂-philic membrane is suitable for commercialization, owing to its facile synthesis strategy and excellent performance.

Table 2 Performance comparison of composite Pebax membranes based on polymer blends.

PBE: poly(2-[3-(2H-benzotriazol-2-yl)-4hydroxyphenyl]ethyl methacrylate)-graft-poly(oxyethylene methacrylate); PGP: poly(glycidyl methacrylate-g-polypropylene glycol); POEM: poly(oxyethylene methacrylate); PEGDME: polyethylene glycol dimethyl ether; POP: porous organic polymer; GO: graphene oxide; PTMEG: poly(tetramethylene ether) glycol; PDMS: polydimethylsiloxane; L-61: Pluronics L-61; g: graft.

In summary, a self-polymerization-confined membrane with ultra-stability was successfully fabricated via a facile one-step method. The resulting membranes exhibited superior CO₂/light gas separation performance compared with the conventional membranes, owing to the high content of the amorphous PEO phase. The robust structure, which was derived from the intricate entanglements between poly-PEGMEA and Pebax chains, endows the membrane with excellent mechanical properties and enables an extremely high loading of PEGMEA (up to 90 wt%). Therefore, the separation performance easily surpasses the conventional upper-bound limit. Moreover, good structural stability allows further increases in CO₂ permeability (up to 832 Barrer), CO₂/N₂ selectivity (up to 63.5), and CO₂/H₂ selectivity (up to 20.8) at a high feed pressure of 20 bar. Finally, during the long-term stability tests, the gas separation performance exhibited negligible attenuation for up to 120 h. The typical as-developed self-polymerized structure will aid the progress of polymer workability toward the practical application of membranes for addressing various energy and environmental issues.

This work was supported by the National Natural Science Foundation of China (21878062, 21878062, and 22111530113), the Natural Science Foundation of Heilongjiang Province for Distinguished Young Scholars (JQ2020B001), Heilongjiang Touyan Team (HITTY20190033), Fundamental Research Funds for the Central Universities, and State Key Laboratory of Urban Water Resource and Environment (Harbin Institute of Technology) (2020DX02).

Bin Zhu, Shanshan He, Yadong Wu, Songwei Li, and Lu Shao declare that they have no conflict of interest or financial conflicts to disclose.

Supplementary data to this article can be found online at https://doi.org/10.1016/j.eng.2022.03.016.