Bifunctionalized NH 2 -UiO-66-CF 3 mixed matrix membranes for gas separation

Abstract

Bifunctionalized NH₂-UiO-66-CF₃ nanofillers were successfully prepared by grafting - CF₃ on NH₂-UiO-66 via aldehyde-amine condensation reaction and applied to the preparation of 6FDA-ODA/NH₂-UiO-66-CF₃ mixed matrix membranes (MMMs). The successful introduction of -NH₂ and -CF₃ groups was verified by FT-IR, XRD, and XPS characterization, and the integrity of the crystal structure was ensured. Due to the presence of hydrogen bonding, NH₂-UiO-66-CF₃ showed good interfacial compatibility with the polymer matrix. Due to the introduction of the - CF₃ group, the MMMs exhibited excellent thermal stability (T_10% > 500°C, T_g = 300°C–305°C), hydrophobicity (contact angle of 92.96–97.71), and mechanical properties (tensile strength of 74–85 MPa). The gas separation performance tests showed that the MMMs exhibited significantly improved gas permeability and selectivity compared to the pure 6FDA-ODA membranes, which was attributed to the fact that the -NH₂ and -CF₃ groups provided more adsorption sites for gases. In particular, the PI/NUF(10) MMMs showed the best performance, with permeabilities of 42.71, 9.22, and 92.79 Barrer for CO₂, O₂, and He, respectively, which were 248%, 213%, and 173% higher than those of the pure membranes. Meanwhile, the selectivities of CO₂/N₂, O₂/N₂ and He/N₂ were improved by 37%, 24%, and 8%, respectively. Close to the Robeson upper limit, demonstrating excellent gas separation potential.

Keywords

polyimides mixed matrix membranes gas separation

Introduction

The extensive use of non-renewable fossil fuels (e.g., coal, petroleum, and natural gas) has resulted in increasingly severe issues of atmospheric pollution and energy scarcity. As a pivotal technology to address these challenges, membrane-based gas separation has gained widespread application in air separation, natural gas purification, and hydrogen recovery owing to its advantages of high efficiency, low energy consumption, and environmental friendliness.^1–5 In membrane separation processes, material selection is critical, with options including inorganic, polymeric, and mixed matrix membranes (MMMs).^6–13 In order to break through the bottleneck of organic polymer membranes constrained by the Robeson upper bound,^14,15 and overcome the shortcomings of difficult and costly membrane fabrication of inorganic materials, researchers have increasingly focused on MMMs incorporating organic polymers and inorganic porous fillers. The introduction of fillers not only disrupts the close packing of polymer chains and increases the free volume, which in turn enhances membrane permeability,^16,17 but also forms interfacial pores and nanocavities in MMMs, providing fast transport channels for gases. MMMs combine the excellent mechanical properties and processability of organic membranes with the high permeability and selectivity of inorganic fillers, demonstrating significant industrial potential. However, poor interfacial compatibility between fillers and polymer matrices remains a major challenge hindering commercial implementation.

Metal-organic frameworks (MOFs) are crystalline porous materials formed by the coordination between metal nodes and organic ligands, exhibiting distinct advantages including high porosity, large specific surface area, and tunable functionality.^18–20 Compared with conventional inorganic fillers (e.g., zeolites, graphene, and carbon nanotubes),^21–23 MOFs demonstrate superior compatibility with polymer matrices attributed to their organic ligands, rendering them ideal candidates for MMMs fabrication. Particularly, UiO-66 with zirconium metal node demonstrates exceptional chemical and thermal stability while maintaining structural integrity under harsh conditions, positioning it as a preferred material for composite membrane fabrication.^24,25 However, the incorporation of crystalline MOFs into polymers often induces interfacial defects due to mismatched phase compatibility. To address this issue, ligand functionalization strategies have been developed.^26,27 Amino-functionalized MOFs serve dual roles: (1) providing alkaline sites for enhanced CO₂ adsorption, and (2) forming hydrogen bonds with polymer polar groups to strengthen interfaces. Similarly, fluorinated MOFs concurrently enhance MMMs hydrophobicity, gas selectivity, and thermal stability.^28,29 Notably, Pu et al.³⁰ reported that the NUS-8-NH₂/PIM-1 MMM achieved a CO₂ permeability of 14,000 Barrer with CO₂/N₂ selectivity of 30, surpassing the 2008 Robeson upper bound. This enhancement originates from -NH₂ groups acting as both CO₂ transport promoters and interfacial linkers via hydrogen bonding with PIM’s cyano groups. Zhou et al.³¹ further demonstrated that UiO-66-(CF₃)₂ incorporation improved both separation performance and hydrothermal stability of PIM-1 membranes. These cases validate that MOFs functionalization effectively enhance membrane performance through tailored interfacial engineering. Nevertheless, single-functionalized MOFs exhibit certain limitations in gas separation performance enhancement. Bifunctional MOFs offer multiple advantages including additional gas adsorption sites, improved selectivity and adsorption capacity for various gas species, as well as enhanced membrane stability under diverse operating conditions. Park et al.³² developed a mixed-ligand strategy to prepare amine- and fluorine-co-functionalized MIL-101(Cr). This bifunctional material exhibited significantly enhanced CO₂ separation performance due to its dual adsorption sites. Remarkably, the co-functionalized MIL-101(Cr) retained excellent separation efficiency even in a humid environment, confirming its practical application potential.

In this work, NH₂-UiO-66 with optimal pore size is first synthesized and the dual-functionalized NH₂-UiO-66-CF₃ filler is subsequently prepared by introducing -CF₃ groups into the MOF through an aldehyde-amine condensation reaction where -CF₃ partially substitutes -NH₂ groups. The modified MOF is then incorporated into the 6FDA-ODA matrix, a benchmark material for gas-separation membranes, to prepare 6FDA-ODA/NH₂-UiO-66-CF₃ MMMs with different MOF loadings. The MMM structures are systematically characterized by FT-IR and XRD, while their thermal stability, hydrophobicity, mechanical properties and gas separation performance are comprehensively evaluated. These analyses reveal the crucial role of NH₂-UiO-66-CF₃ bifunctionalization in enhancing membrane characteristics, particularly gas separation efficiency, and clarify the underlying mechanisms.

Experimental

Materials

4,4′-(hexafluoroisopropylidene)diphthalic anhydride (6FDA, purity > 98%, Adamas), 4,4′-oxydianiline (ODA, 98%, Adamas), m-Cresol (99%, Shanghai Aladdin Co., Ltd), isoquinoline (97%, Shanghai Aladdin Co., Ltd), Zirconium (IV) chloride (ZrCl₄, purity > 98%, Adamas), 2-Aminoterephthalic Acid (BDC-NH₂, purity >98%, Adamas), 4-(Trifluoromethyl)Benzaldehyde (4-TFMB, purity > 98%, Adamas), N,N-dimethylformamide (DMF, Sinopharm Chemical Reagent Co., Ltd). Acetic acid (CH₃COOH) and ethanol (EtOH), methanol (MeOH) (Xilong Chemical Co., Ltd). All solvents are analytical grade and can be used without further treatment. CO₂, O₂, N₂ and He (Liuzhou Hongrun Gas Co., Ltd.) are experimental gases with purity greater than 99.9%.

Preparation of 6FDA-ODA

6FDA (4.44 g, 0.01 mol) and ODA (2.00 g, 0.01 mol) at a 1:1 molar ratio are placed in a 100.0 mL three-necked flask equipped with mechanical stirring and a condenser. M-Cresol (58.0 mL) is added, and the reaction system is heated to 50°C under nitrogen atmosphere. After complete mixing, isoquinoline (0.3 mL) is added as the catalyst. The thermal imidization proceeds via stepwise heating: first at 80°C for 3 h, then at 120°C for 3 h, followed by 180°C for 3 h, and finally at 200°C for 12 h. Upon reaction completion, the system is cooled to room temperature, the condenser water and nitrogen flow are stopped, and the solution is poured slowly into methanol for precipitation. The precipitated product is collected by filtration and dried under vacuum at 150°C for 12 h to yield fibrous polyimide.

Synthesis of MOFs

The synthesis of NH₂-UiO-66 powder followed a literature method with modifications.³³ First, ZrCl₄ (1 mmol) and BDC-NH₂ (1 mmol) were dispersed in a mixed solvent (20.0 mL of DMF and 5.0 mL of acetic acid, total volume 25.0 mL) via ultrasonication, followed by stirring until a homogeneous mixture was obtained. The solution was then transferred to a Teflon-lined autoclave and heated at 120°C for 24 h. After cooling to room temperature, the solid product was collected by centrifugation and washed repeatedly with DMF and methanol. Finally, the powder was dried at 120°C for 24 h to yield NH₂-UiO-66.

NH₂-UiO-66-CF₃ was synthesized via post-synthesis modification (PSM)³⁴ as follows: first, NH₂-UiO-66 was activated under vacuum at 150°C for 24 h to remove residual solvent molecules from the pores. Subsequently, 0.40 g of the activated NH₂-UiO-66 was dispersed in 60.0 mL of ethanol via ultrasonication and stirred until homogeneous. Then, 3.0 mL of 4-trifluoromethylbenzaldehyde (4-TFMB) was added to the mixture, and the reaction was refluxed at 80°C for 12 h under nitrogen protection. After cooling to room temperature, the solid product was collected by centrifugation and washed thoroughly with methanol several times to remove unreacted reagents. Finally, the product was dried under vacuum at 80°C for 24 h to yield NH₂-UiO-66-CF₃ powder.

Fabrication of membranes

The gas separation membranes were fabricated by a solution-casting method (Scheme 1). Five membrane compositions with filler loadings of 0, 5, 8, 10 and 15 wt% were prepared. The detailed procedure was as follows: first, the NH₂-UiO-66-CF₃ filler was ultrasonically dispersed in DMF, and then the filler dispersion was added into the 6FDA-ODA solution dissolved in DMF, followed by prolonged ultrasonic stirring to ensure homogeneous dispersion. The resulting mixture was doctor-bladed onto a leveled glass substrate, with solvent evaporation at 80°C. After cooling to room temperature, the membrane was peeled off using lukewarm water. Finally, the membrane was vacuum-dried at 150°C for 12 h. The mixed matrix membranes are hereafter denoted as PI/NUF(x), where x represents the filler loading (calculated using equation (1)).

f i l l e r l o a d i n g (w t %) = \frac{f i l l e r w e i g h t}{f i l l e r w e i g h t (g) + p o l y m e r w e i g h t (g)} \times 100 %

(1)

Characterizations

The crystalline properties of MOF fillers and membranes were characterized by X-ray diffraction (XRD, X’Pert PRO X, PANalytical B.V.). Fourier transform infrared spectrometry (FT-IR, Nicolette 6700-NXR, Thermo Nicolet Corporation) and X-ray photoelectron spectroscopy (XPS, Escalab 250Xi, USA Thermo Electron Corporation) were used to characterize the MOF fillers and membrane structural properties. Thermal properties were evaluated through thermogravimetric analysis (TGA, STA 449 C, NETZSCH) and differential scanning calorimetry (DSC, DSC 214, NETZSCH) under N₂ atmosphere (10°C/min heating rate, TGA: RT-800°C, DSC: 30°C–400°C). The MOF filler was vacuum degassed at 150°C for 12 h, and then the N₂ adsorption-desorption isotherms of the filler were measured by a specific surface area and pore size analyzer (TriStar II 3020, Micromeritics). The specific surface area of the sample was calculated using the Brunauer-Emmett-Teller (BET) method. The micromorphology of MOF powder as well as membrane was observed by scanning electron microscopy (SEM, Hitachi S-4800, Hitachi). Surface hydrophobicity was assessed through water contact angle measurements (JY PHb, Chengde Youte Instrument Co.). An electronic universal testing machine (CMT2103, MTS Corporation) with a maximum tensile load of 5 kN was used to test the mechanical properties of the films.

Pure gas separation evaluation

The gas permeability coefficients (P) of pure CO₂, O₂, He, and N₂ were measured by the constant volume/variable pressure method (Labthink VAC-V1, Jinan, China) at 35°C with 0.4 MPa feed pressure. Prior to testing, the dried membrane was mounted between the test chambers. After tightening the upper chamber, the system underwent sequential evacuation: (1) lower chamber evacuation for 60 s, followed by (2) full-system evacuation for ≥3 h to achieve required vacuum levels. The upper chamber was then pressurized with test gas to establish a constant transmembrane pressure difference, driving gas permeation from the high- to low-pressure side. Triplicate measurements were conducted for each membrane under identical conditions to ensure data reproducibility, with results averaged to minimize experimental error.

Permeability (P) is calculated as follows:

P = \frac{V_{d} l}{{A T P}_{u p} \times 0.278} \times \frac{d_{p}}{d_{t}} \times 10^{10}

(2)

where, P is the permeability coefficient (Barrer, 1 Barrer = 1 × 10⁻¹⁰ cm³ [STP] cm/(cm² s cm Hg)), V_d is the calibrated permeation volume (cm³), l and A are the membrane thickness (μm) and the effective test area (cm²), respectively, T(K) is the operating temperature, P_up is the upstream pressure (cm Hg), and dp/dt is the steady-state pressure growth rate.

The ideal selectivity (α_A/B) of a pair of gases is the ratio of their permeability coefficients, which can be expressed as:

α_{A / B} = \frac{P_{A}}{P_{B}}

(3)

P_A and P_B are the permeability coefficients of pure gases A and B, respectively.

Results and discussion

Synthesis and structural characterization of NH₂-UiO-66-CF₃

Figure 1 illustrated the XRD patterns of NH₂-UiO-66 and NH₂-UiO-66-CF₃. Both samples exhibited diffraction profiles matching the simulated patterns,³⁵ and prominent characteristic peaks appeared at the (111) and (002) crystal planes (2θ = 7.4°, 8.5°), indicating that the crystal structure of NH₂-UiO-66 was successfully synthesized. After trifluoromethylation, NH₂-UiO-66-CF₃ maintained the parent topology while preserving high crystallinity. SEM characterization (Figure S1) revealed well-defined octahedral nanoparticles (average size: 200 nm) with sharp edges and intact surfaces, demonstrating structural preservation during trifluoromethylation.³⁶ The high consistency of this topological morphology combined with XRD data conclusively demonstrated that the trifluoromethyl modification process did not lead to the reconstruction of the MOF crystal structure, providing a material basis for the subsequent interface compatibility study.

In the FTIR spectrum (Figure S2), NH₂-UiO-66 and NH₂-UiO-66-CF₃ exhibited asymmetric and symmetrical stretching vibration peaks of - C = O at 1690 cm⁻¹ and 1380 cm⁻¹, as well as longitudinal and transverse vibration peaks of Zr-O bonds at 768 cm⁻¹ and 669 cm⁻¹. Notably, in the NH₂-UiO-66-CF₃ curve, the -NH₂ characteristic absorption peak near 3500 cm^-1 was significantly weakened, and a new -CF₃ characteristic absorption peak appeared at 1060 cm⁻¹, which indicated that some -NH₂ groups were replaced by -CF₃. In addition, the full spectrum of NH₂-UiO-66-CF₃ and the fine spectrum of elements were measured by XPS to further determine its chemical structure composition. It can be seen from the wide-scan XPS spectrum of Figure 2(a) that NH₂-UiO-66-CF₃ is composed of five elements: C, N, O, F and Zr, which also indicated that -CF₃ was successfully grafted onto NH₂-UiO-66. Figure 2(b) showed that the high-resolution XPS spectra of C 1s were convoluted into C-C (284.40 eV), C=N (285.76 eV), O-C=O (288.83 eV) and C-F (292.96 eV) bonds. The N 1s spectrum (Figure 2(c)) showed characteristic peaks for -NH₂ (399.43 eV) and C=N (400.43 eV). The presence of the C-F bond was further confirmed by the F 1s peak at 688.07 eV (Figure 2(d)). All experimental results showed that 4-TFMB was successfully grafted onto NH₂-UiO-66 through aldehyde amine condensation reaction, while the -NH₂ group did not completely disappear, and finally the bifunctional NH₂-UiO-66-CF₃ was obtained.

The N₂ adsorption and desorption curves of both NH₂-UiO-66 and NH₂-UiO-66-CF₃ (Figure 3) exhibited type I and type IV isotherms, indicating that hierarchical porous MOF materials with co-existing microporous and mesoporous sites were successfully prepared by the acetic acid modifier. The presence of defects led to the exposure of more unsaturated sites, which enhanced its potential application in gas separation. The BET specific surface area of NH₂-UiO-66 was 648 m² g⁻¹ (Table S1), representing a significant increase compared to the literature sample without acetic acid (575 m² g⁻¹).³⁷ This improvement was mainly attributed to the strategy of adjusting the concentration of acetic acid to control the particle size, which effectively increased the specific surface area of the MOF. The slight decrease in the specific surface area and microporous volume of NH₂-UiO-66-CF₃ compared to NH₂-UiO-66 was attributed to the modification of 4-TFMB that resulted in the clogging of some of the pores. However, F atoms were shown to enhance host-guest interactions, leading to significantly higher adsorption capacities for CO₂ and CH₄ than NH₂-UiO-66.³³ Additionally, the reduced pore volume (V_Micro = 0.23 cm³ g⁻¹) of NH₂-UiO-66-CF₃ created pore channels more favorable for He (2.60 Å) diffusion, making it a highly promising porous material for He separation.

The effect of functionalization modifications on the thermal stability of the materials was systematically evaluated by thermogravimetric analysis (TGA) (Figure 4). Both MOF fillers, NH₂-UiO-66 and NH₂-UiO-66-CF₃, exhibited a three-stage weight loss: the first stage (30°C–150°C) corresponded to the loss of free water molecules, the second stage was attributed to the removal of solvent guest molecules (e.g., DMF) and bond energy changes induced by - NH₂/-CF₃ group, and the final stage at about 350°C for the weight loss due to the combustion of 2-aminoterephthalic acid and decomposition of the MOF structure.³⁸ NH₂-UiO-66-CF₃ exhibited a significantly enhanced thermal stability, with a higher decomposition onset temperature and a slower weight loss rate during framework decomposition compared to the parent material. This enhanced effectivity of thermal stability was attributed to the strong electronegativity of trifluoromethyl groups in synergy with Zr-O bonds. The difference in the second weight loss region further confirmed the successful grafting of -CF₃ groups, providing thermodynamic stability for composite membranes under high-temperature conditions.

Synthesis and characterization of PI/NH₂-UiO-66-CF₃ MMMs

The membrane structure was analyzed by XRD pattern (Figure 5). The pure 6FDA-ODA matrix membrane showed a typical amorphous broadening diffraction peak at 2θ = 15°, while all PI/NUF(x) MMMs clearly displayed sharp characteristic peaks of the MOF fillers at 7.4°(111) and 8.5°(002) while retaining this feature.³⁹ As the NH₂-UiO-66-CF₃ loading increased from 5 wt% to 15 wt%, the characteristic peak intensity of MOF increased, confirming the successful incorporation of functionalized UiO-66 fillers without crystal structure degradation during MMMs preparation. The FT-IR spectra of the membranes were shown in Figure 6, with asymmetric and symmetric stretching vibration peaks of C=O at 1784 cm⁻¹ and 1727 cm⁻¹, stretching vibration peaks of C-N, C-O, and C-F at 1380 cm⁻¹, 1154 cm⁻¹ and 1113 cm⁻¹, respectively, and a bending vibration peak of the imide ring at 723 cm⁻¹, which confirmed the successful synthesis of high-temperature polyimide after thermoimination.^40,41 Notably, the C=O vibrational peak of 1784 cm⁻¹ gradually weakened with higher filler content, likely due to hydrogen bonding between -NH₂ of MOF nanoparticles and C=O of PI.⁴² Additionally, the appearance of a Zr-O stretching vibration at 660 cm⁻¹, with intensity proportional to filler content, further evidenced NH₂-UiO-66-CF₃ incorporation into the polymer.

The dispersion of MOF in polymer solution greatly influenced the gas separation performance of MMMs. SEM analysis of 6FDA-ODA and composite membrane cross-sections at different loadings revealed that all MMMs exhibited a concave morphology (Figure 7), characteristic of PI polymer chain accumulation. Octahedral MOF particles were dispersed in the MMMs, and the number of dispersed particles increased with higher loadings of NH₂-UiO-66-CF₃. However, agglomeration became noticeable at elevated loadings, slightly reducing filler-polymer affinity. Despite the progressively rougher cross-sectional morphology, no interfacial defects were observed, which was attributed to the fact that the presence of bifunctional groups in NH₂-UiO-66-CF₃ enhanced the interaction with PI, thereby improving the interfacial compatibility.

The thermal stability of 6FDA-ODA and PI/NUF(x) membranes was characterized by TGA, as illustrated in Figure 8. The TGA curves of the composite and matrix membranes displayed similar two-stage degradation: the loss of solvent and decomposition of NH₂-UiO-66-CF₃ up to 500°C, and the decomposition and the carbonization of the PI in the second phase. The maximum weight loss rates of all MMMs occurred above 500°C and the residual carbon rates were higher than 50% (Table S2), indicating that the MMMs have good thermal stability. However, PI/NUF(15) showed a slightly lower char yield (52%) than the matrix membrane (55%), attributed to the aggregation of NH₂-UiO-66-CF₃ in the polymer membrane at higher filler loadings, potentially weakening the filler-PI interactions.⁴³ This phenomenon was further verified by DSC test. As shown in Figure 9, the glass transition temperature (T_g) varied between 299°C and 305°C, first increasing and then decreasing. All MMMs had only one T_g, while the addition of fillers limited the movement of polymer chains, resulting in higher T_g for all MMMs than pure membranes.⁴⁴ Although PI/NUF (15) showed a decrease in T_g due to the weakening of polymer-filler interactions caused by filler agglomeration,⁴⁵ there was no phase separation, indicating good interfacial compatibility between NH₂-UiO-66-CF₃ and PI.

The water contact angle test (θ_w) was performed on the pure membrane 6FDA-ODA and its MMMs to evaluate hydrophobicity, where larger θ_w values indicated greater membrane hydrophobicity. As shown in Figure 10, 6FDA-ODA membrane (θ_w = 98.85°) exhibited inherent hydrophobicity due to its -(CF₃)₂ groups.⁴⁶ With the increase of the bifunctionalized NH₂-UiO-66-CF₃ loading, θ_w showed a gradual decrease, but the decrease slowed down significantly, only slightly from 97.71 to 92.96. This was attributed to the strong hydrophobic effect of -CF₃ in the MOF filler partially offsetting the hydrophilicity of -NH₂. However, the hydrophobic enhancement effect was somewhat limited by the insufficient substitution of the -CF₃ group. Notably, all PI/NUF(x) membranes maintained θ_w > 90°, demonstrating superior water-resistant properties for gas separation applications. This characteristic effectively minimized water molecule interference, thereby enhancing gas separation efficiency and long-term stability.

The mechanical properties of the 6FDA-ODA matrix membrane and its mixed matrix membranes (MMMs) were evaluated through tensile testing (Figure 11). The tensile strengths of all membranes were in the range of 74–94 MPa, the elongation at break was in the range of 6.07%–7.24%, and the Young’s modulus was in the range of 1.54–1.91 GPa (Table S3). The MMMs showed slightly reduced tensile strength and elongation at break compared to the pure 6FDA-ODA membrane, primarily due to the incorporation of rigid NH₂-UiO-66-CF₃ filler that both restricted polymer chain mobility and disrupted chain regularity. However, the Young’s modulus of the MMMs increased compared to the 6FDA-ODA membrane, because of the formation of strong hydrogen bonds between - NH₂ and the polymer, as well as hydrophobic interactions, dipole-dipole interactions and van der Waals forces between -CF₃ and 6FDA-ODA. These intermolecular interactions restricted polymer chain mobility, resulting in enhanced film rigidity.^47,48 Among them, PI/NUF(10) membrane exhibited the largest Young’s modulus, indicating that the MMM containing 10 wt% NH₂-UiO-66-CF₃ had the most excellent interfacial interaction force.

Gas separation performance

The permeability of the prepared membranes to four pure gases, CO₂, O₂, He, and N₂, was measured using the constant volume-variable pressure method at 35°C and 0.4 MPa feed pressure. The trend of gas permeability versus selectivity was observed in Figure 12. He had the largest permeability among the four gases, because it had the smallest molecular size (2.60 Å) and the highest diffusion coefficient, which facilitated transport through the membrane’s microporous structure. With the increase of NH₂-UiO-66-CF₃ content, the gas permeability gradually improved, mainly attributed to the MOF filler introducing additional pores that enhanced gas diffusion while simultaneously increasing the membrane’s free volume through disruption of molecular chain stacking, thereby resulting in smoother gas transport.⁴⁹ The permeation performance of the composite membranes reached a maximum when the loading of NH₂-UiO-66-CF₃ reached 10 wt%, with permeabilities of 42.71 Barrer for CO₂, 9.22 Barrer for O₂, 92.79 Barrer for He, and 1.77 Barrer for N₂ (Table 1). Compared to the pure matrix membrane 6FDA-ODA, the permeability coefficients of CO₂ were enhanced by 248%, O₂ by 213%, and He by 173%. Meanwhile the selectivities of CO₂/N₂, O₂/N₂ and He/N₂ improved by 37%, 24% and 8%, respectively. This indicated that the appropriate amount of MOF packing could effectively enhance the permeability and selectivity of the membranes. However, when the filler loading was further increased to 15 wt%, the gas permeability showed a decreasing trend, which was attributed to the possible aggregation of excess NH₂-UiO-66-CF₃ in the membranes, leading to the blockage of some of the pores or the formation of disordered filler stacking, which in turn reduced the gas diffusion rate.⁵⁰ Nevertheless, the permeation rates of the PI/NUF (15) MMM still maintained significantly higher than the 6FDA-ODA matrix membrane, demonstrating that the composite membranes could preserve better permeability performance even at higher loadings. It was noteworthy that the gas selectivity of MMMs began to show a declining trend before the filler content reached 15 wt%. This phenomenon may be attributed to nanoparticle agglomeration during the filler loading process, resulting in uneven dispersion within the membrane. Such agglomeration could potentially create non-selective voids that preferentially increase N₂ permeation, thereby reducing overall selectivity.

Table 1.

Comparison of the gas separation properties of MMMs reported in this work with those reported in the literature.

Membranes	Permeability (Barrer^a)				Ideal selectivity α^b
Membranes	CO₂	O₂	He	N₂	CO₂/N₂	O₂/N₂	He/N₂	He/CO₂
6FDA-ODA	12.29	2.95	34.02	0.77	17.56	4.21	48.60	2.77
PI/NUF(5)	18.12	4.59	41.94	0.79	22.94	5.81	53.09	2.31
PI/NUF(8)	32.44	6.97	67.25	1.23	26.37	5.67	54.67	2.07
PI/NUF(10)	42.71	9.22	92.79	1.77	24.13	5.21	52.42	2.17
PI/NUF(15)	36.44	7.71	78.34	1.75	20.82	4.41	44.77	2.15
6FDA-ODA/NH₂-MIL-53 (20)-A⁴⁷	35.2	—	—	—	14.6	—	—	—
PI-tb???o-66-NH₂(30)⁵²	415	92.3	—	22.0	18.9	4.2	—	—
UIO-66⁵³	37.60	—	—	1.76	21.40	—	—	—
UIO-66-5⁵⁴	298.2	—	—	16.38	18.2	—	—	—
ZIF-7-I-PI2(15)⁵⁵	—	2.9	—	0.19	—	15.2	—	—
PSf/ZIF-8 (16)⁵⁶	28.4	16.2	61.0	15.0	1.89	1.08	4.07	2.57
Matrimid 5218/mil-101 (10)⁵⁷	6.95	—	—	0.131	53	—	—	—
PSF/Silica (20)⁵⁸	19.70	5.00	27.70	1.12	17.59	4.46	24.73	1.41
TR-PBOI/CNTs⁵⁹	126	—	—	9.3	13.60	—	—	—
PU/4A Zeolite (18)⁶⁰	93.28	9.76	—	3.85	24.22	2.53	—	—

Note. 1 Barrer = 10⁻¹⁰ [cm³(STP)·cm]/(cm²s·cmHg), Ideal Selectivity α = P(A)/P(B), A and B are two different pure gases.

In addition, the introduction of NH₂-UiO-66-CF₃ significantly enhanced CO₂ permeability and CO₂/N₂ selectivity. This enhancement primarily resulted from strong interactions between CO₂ molecules and polar groups (-NH₂ and -CF₃) in the MOF fillers. The N, O, and F elements provided alkaline active sites for CO₂ adsorption, while C-F and N-H bonds generated dipole-quadrupole interactions with CO₂ molecules, substantially increasing CO₂ solubility in the membrane.^38,51 Thus, despite CO₂ having a larger kinetic diameter than He, the MMMs with 10 wt% MOF loading demonstrated optimal enhancement in both CO₂ permeability and CO₂/N₂ selectivity compared to the pristine 6FDA-ODA matrix membrane, showing remarkable improvement of 248% and 37%, respectively. By optimizing the NH₂-UiO-66-CF₃ loading, the composite membranes achieved not only enhanced overall gas separation performance but also superior CO₂ separation efficiency while maintaining satisfactory mechanical properties, demonstrating great potential for gas separation applications.

Figure 13 demonstrated the gas separation performance of PI/NUF(x) MMMs in comparison with other reported mixed matrix membranes relative to the Robeson upper bound. The results showed that incorporating NH₂-UiO-66-CF₃ into 6FDA-ODA significantly enhanced both gas permeability and selectivity, with performance gradually approaching the Robeson upper bound. Among the tested membranes, PI/NUF(10) membrane exhibited the most prominent performance for CO₂/N₂, He/N₂, and He/CO₂ gas pairs, closely approaching the upper bound, while PI/NUF(8) membrane is infinitely close to the upper bound for O₂/N₂ separation. These performance variations resulted from differences in gas physicochemical properties, including molecular size, solubility parameters, and diffusion coefficients. Compared with literature reports (Table 1), this study successfully overcame the traditional trade-off effect through bifunctionalized NH₂-UiO-66-CF₃ incorporation, achieving simultaneous enhancement in both permeability and selectivity. These findings demonstrated that bifunctionalized MOF fillers represented an effective strategy for optimizing MMMs performance and showed significant potential for industrial gas separation applications.

Conclusion

In this study, hierarchical porous NH₂-UiO-66 with defects was first synthesized using an acetic acid modifier, followed by successful preparation of bifunctionalized NH₂-UiO-66-CF₃ nanoparticles through -CF₃ grafting via an aldehyde-amine condensation reaction, achieving simultaneous coexistence of -NH₂ and -CF₃ groups. This bifunctional design conferred multiple advantages to the materials, as demonstrated through thermogravimetric analysis, contact angle measurements, and tensile tests, which confirmed that the PI/NUF(x) MMMs exhibited good thermal stability, hydrophobicity and mechanical properties. The strong hydrogen bonding between -NH₂ groups and 6FDA-ODA significantly enhanced the filler-matrix interfacial affinity, while the hierarchical pore structure effectively strengthened the molecular sieving effect. Furthermore, the combined effects of (1) dipole-quadrupole interactions between N-H/C-F bonds and CO₂ molecules, and (2) Lewis basic active sites provided by N and F atoms, synergistically enhanced both permeability and selectivity for CO₂, O₂, and He in PI/NUF(x) MMMs, overcoming the traditional permeability-selectivity trade-off limitation. Through this “dual adsorption sites with hierarchical pores” synergistic mechanism, the bifunctionalized NH₂-UiO-66-CF₃ demonstrated significant potential for separation applications, offering an efficient new approach for carbon capture from flue gas and air separation processes.

Supplemental Material

Supplemental Material - Bifunctionalized NH₂-UiO-66-CF₃ mixed matrix membranes for gas separation

Supplemental Material for Bifunctionalized NH₂-UiO-66-CF₃ mixed matrix membranes for gas separation by Zhiqing Liang, Zhongying Liu, Xiuyuan Wei, Hui Wu, Chanjuan Liu and Xiaohua Huang in High Performance Polymers.

Footnotes

ORCID iDs

Chanjuan Liu

Xiaohua Huang

Funding

The authors disclosed receipt of the following financial support for the research: This work was supported by the National Natural Science Foundation of China (51963007), Guangxi Science and Technology Plan Project (GuikeAD25069100), Guangxi Natural Science Foundation (2021GXNSFAA075016, 2020GXNSFAA159009, 2020GXNSFAA159006), Independent Research Project of Guangxi Colleges and Universities Key Laboratory of Natural and Biomedical Polymer Materials (24POLY-AA-02), Opening Fund of Key Laboratory of New Processing Technology for Nonferrous Metal & Materials, Ministry of Education/Guangxi Key Laboratory of Optical and Electronic Materials and Devices (No. 22KF8).

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.*

Supplemental Material

Supplemental material for this article is available online.

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Supplementary Material

Please find the following supplemental material available below.

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Bifunctionalized NH 2 -UiO-66-CF 3 mixed matrix membranes for gas separation

Abstract

Keywords

Introduction

Experimental

Materials

Preparation of 6FDA-ODA

Synthesis of MOFs

Fabrication of membranes

Characterizations

Pure gas separation evaluation

Results and discussion

Synthesis and structural characterization of NH2-UiO-66-CF3

Synthesis and characterization of PI/NH2-UiO-66-CF3 MMMs

Gas separation performance

Conclusion

Supplemental Material

Supplemental Material - Bifunctionalized NH2-UiO-66-CF3 mixed matrix membranes for gas separation

Footnotes

ORCID iDs

Funding

Declaration of conflicting interests

Data Availability Statement

Supplemental Material

References

Supplementary Material

Synthesis and structural characterization of NH₂-UiO-66-CF₃

Synthesis and characterization of PI/NH₂-UiO-66-CF₃ MMMs

Supplemental Material - Bifunctionalized NH₂-UiO-66-CF₃ mixed matrix membranes for gas separation