WO2022173375A1

WO2022173375A1 - Membrane composites, methods and uses thereof

Info

Publication number: WO2022173375A1
Application number: PCT/SG2022/050066
Authority: WO
Inventors: Bofan LI; Ji Wu; Susilo JAPIP; Kai Yu Wang; Tai-Shung Chung
Original assignee: National University Of Singapore
Priority date: 2021-02-10
Filing date: 2022-02-09
Publication date: 2022-08-18

Abstract

This disclosure concerns membrane composites, the methods of fabricating the membrane composites and their uses. The membrane composite comprises a layer of polymer and a macrocycle homogenously distributed within the polymer layer. The polymer comprises a first polar moiety and the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer. In some embodiments, calixarene is used to be infiltrated into a polymer scaffold of an amidoxime-functionalized intrinsic microporosity PIM-1 (AOPIM1), polyamide, or a crosslinked polyimide-amide polymer or polybenzimidazole.

Description

Membrane Composites, Methods and Uses Thereof

Technical Field

The present invention relates, in general terms, to membrane composites, the methods of fabricating the membrane composites and their uses thereof.

Background

Polymer membranes present a promising energy-efficient technology for many important gas separation processes, such as hydrogen recovery (H2/N2, H2/CH4) and carbon capture from syngas, coal gasifiers (H2/CO2) or natural gas (CO2/CH4), which are all critical for clean energy and environmental development. However, the natural random packing of polymers causes broadly distributed pore sizes in these membranes, which unfortunately include non-size-selective microporous regions that could impair their gas separation efficiency and hence industrial competitiveness. The facile control of such non-selective regions with the required size-sensitivity on the sub-angstrom level still remains a key material challenge.

Utilizing the intrinsic molecular-sieving nanopores of an external nanoporous agent, like 2D or 3D inorganic, carbon-based or extended framework-type materials, in polymer gas separation membranes has attracted broad research interest because of their potential promise for realizing advantageous composite properties. However, their poor solution processability and challenging compatibility with organic polymers often lead to detrimental issues of particle agglomeration, interfacial voids or pore blockage which conventionally require delicate interfacial design and more aggressive physical mixing or dispersing techniques to be alleviated. Even for purely organic nanoporous fillers with better polymer affinity and organic-solvent processability, such as covalent organic frameworks (COFs) or porous organic cages (POCs), true molecular mixing or dispersion between individual filler molecules and polymer chains remains difficult due to their inherent tendency to crystallize and self-assemble which perpetuates the threat of nanoscopic interfacial defects. In this respect, for most of the studied filler choices, only a conventional mixed-matrix structure comprising a discretely dispersed filler phase and a continuous polymer phase has been achieved (Figure 1). Besides, physical mixing typically carried out before membrane fabrication often creates an uncontrollable window for this crystallization/self-assembly propensity to be displayed as well as for the polymer chains to block the fillers' nanopores, partly due to the high migration and chain mobility in the solution-mixture state.

Further, organic solvents are intensively employed in pharmaceutical syntheses. A green and efficient practice to recycle waste organic solvents instead of incineration is imperative for earth sustainability. However, traditional separation processes, such as distillation and evaporation, suffer from high energy consumption, large footprints, and environmental unfriendliness. Membrane-based separation to recycle waste organic solvents is emerging in recent decades by means of hydraulic or osmotic pressure as the driving force. Among them, organic solvent nanofiltration (OSN) is widely studied but the high operating pressure and fouling tendency may increase the capital and operating costs. Recently, organic solvent forward osmosis (OSFO) is found to be promising to lower the fouling tendency and treat the highly concentrated feed solutions. Similar to forward osmosis (FO) in water reuse processes, both valuable products and organic solvents can be recovered from pharmaceutical streams. A membrane with both high permeability and selectivity is desirable to yield good OSFO performance. Nevertheless, little work has been done in the field of OSFO, especially for the development of OSFO membranes.

Hydration Technology Innovations (HTI) was the pioneer in developing FO membranes from cellulose acetate (CA), where the selective and support layers were formed integrally via phase inversion. Afterward, thin-film composite (TFC) membranes for seawater reverse osmosis (RO) were modified and adopted for the fabrication of FO membranes. The TFC configuration was superior to the asymmetric, integrally skinned, one in HTI CA membranes because its selective layer and porous substrate could be fabricated separately. As a result, TFC FO membranes generally have a higher flux and a lower internal concentration polarization (ICP) than HTI CA membranes. Thin-film nanocomposite (TFN) FO membranes were further developed to improve the performance. Nano-fillers, such as carbon nanotubes, graphene oxides, metal oxide nanoparticles, and metal-organic framework, have been incorporated into the polyamide selective layers to provide extra free volumes for solvent transport. However, the weak affinity between nano-fillers and the polyamide network may lead to particle aggregation and sacrifice the selectivity. In addition, the construction of transporting channels by nano-fillers usually compromise the selectivity because of their interference to form a dense polyamide network.

As mentioned, organic solvent nanofiltration (OSN), also known as solvent resistant nanofiltration, is a membrane-based separation process dealing with solutes ranging from 200 Da to 1000 Da in organic solvents. It has been reported that organic solvents account for 80%-90% mass utilization in pharmaceutical and fine chemical industries and they are incinerated after one-time use. OSN is designed to recycle and reuse these organic solvents as well as to concentrate and recover valuable products. Thus, it is emerging as a sustainable technology to replace the traditional separation processes, such as distillation and crystallization, or to hybrid with them to reduce the overall energy consumption, operational cost and environmental impact.

Currently, most OSN membranes are produced in a flat-sheet configuration due to its ease of fabrication and modifications. Meanwhile, hollow fiber membranes (HFMs) remain scarce although they have several advantages, such as a large surface to volume ratio, a small footprint and a self-supporting characteristic. Up to now, there is no commercial OSN HFM and only a few lab-made ones are developed in the literature. The commonly used polymers to fabricate OSN FIFMs are polyimides (e.g. P84, Torlon^®), polyacrylonitrile (PAN) and polybenzimidazole (PBI), because of their superior chemical and mechanical stability. Similar to the manufacture of FIFMs for gas and water applications, the dry-jet wet spinning technique is mostly adopted where water is usually used as the coagulant. The as-spun FIFMs can be directly used either as integrally skinned asymmetric membranes or substrates for the fabrication of thin-film composite (TFC) membranes. Although TFC membranes tend to be superior to integrally skinned asymmetric membranes in terms of selectivity for OSN, more steps and cautiousness are required to produce TFC FIFMs. In addition, both substrate and selective layer may experience different degrees of swelling in organic solvents, resulting in defects and loss in selectivity. Thus, integrally skinned asymmetric membranes are more structurally favorable for OSN FIFMs; however, their relatively low selectivity remains a drawback to be solved.

It would be desirable to overcome or ameliorate at least one of the above-described problems.

Summary The present invention is predicated on the understanding that, if a nanoporous agent could be incorporated into polymer membranes on a molecularly homogeneous basis after the polymers are fabricated into films or selective layers, the need of undergoing any potentially problematic prefabrication mixing stages could be eliminated (i.e. the direct insertion of entire fillers rather than via the pre-seeding process or in situ syntheses). In particular, it was found that macrocycles are advantageous in acting as pore regulators in membrane composites. The present inventions provides a membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; and wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer.

In some embodiments, the first polar moiety resides in a polymer backbone, as a side group, or a combination thereof. In some embodiments, the first polar moiety is selected from cyano, acyl, oxyacyl, acyloxy, amino, acylamino, aminoacyl, amidoximyl, oximyl, hydrazonyl, iminyl, hydroxyl, pyrrolyl, oxazolyl, imidazolyl, pyrazolyl, triazolyl, and a combination thereof.

In some embodiments, the macrocycle or derivative thereof is selected from cyclodextrin, calixarene, cucurbituril, resorcinarene, pillararene, and a combination thereof.

In some embodiments, the macrocycle or derivative thereof is a calixarene or a derivative thereof, selected from sulfocalixarene (or sulfonylcalixarene), sulfothiacalixarene, carboxylatocalixarene, aminocalixrene, p-phosphonic acid calixarene, or a combination thereof.

In some embodiments, the macrocycle or derivative thereof has monomer residues of 4 to 12. In some embodiments, the macrocycle or derivative thereof has an upper rim and a lower rim, the upper rim and lower rim are separated by a frustum-shaped cavity, wherein a diameter of the lower rim is at least about 3 A.

In some embodiments, the upper rim is functionalised with at least one moiety independently selected from sulfonyl, phosphoryl, amino, carboxyl, oxyalkyl and a combination thereof.

In some embodiments, the lower rim is functionalised with at least one moiety independently selected from optionally substituted acyloxy, optionally substituted acyl, optionally substituted alkyl, optionally substituted alkenyl, optionally substituted alkoxy, and optionally substituted alkenyloxy.

In some embodiments, a mole ratio of macrocycle to polymer is about 0.1% to about 10%.

In some embodiments, the macrocycle is about 0.1 wt% to about 10 wt% relative to the membrane composite.

In some embodiments, the macrocycle is homogenously dispersed across the thickness of the layer of polymer.

In some embodiments, the macrocycle is further covalently bonded to the layer of polymer.

In some embodiments, the polymer in the layer of polymer is selected from a polymer with intrinsic microporosity-1 (PIM-1), polyamide, polyamide-imide, polybenzimidazole, polylactic acid (PLA), polycarboxylic acid (PCA), polyethylenimine (PEI), polyarylamine, polyalkylamine, polyallylamine, poly(vinyl amine) or a combination thereof.

In some embodiments, the polymer layer of polymer is crosslinked.

The present invention also provides a method of fabricating a membrane composite, the membrane composite having a macrocycle homogenously distributed within a layer of polymer, the method comprising incubating the layer of polymer in a macrocycle solution; wherein the polymer comprises a first polar moiety; and wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer.

In some embodiments, the electrostatic interaction is selected from hydrogen-bond, proton transfer, ionic interaction or a combination thereof.

In some embodiments, the macrocycle solution comprises a protic solvent.

In some embodiments, the protic solvent is selected from methanol, ethanol, isopropanol, n-propanol, water, and a combination thereof.

The present invention also provide a membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from acyl, amino, acylamino, aminoacyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 5 to 8.

Advantageously, when these macrocycles are used, it was found that the membrane composite are especially beneficial for use in organic solvent filtration.

In some embodiments, the layer of polymer is crosslinked with a crosslinker having at least 2 pendant groups, the pendant group selected from aminoalkyl, haloalkyl, and a combination thereof. Advantageously, the crosslinker provides addition polar moieties which can be incorporated into the layer of polymer and forms a portion (if not all) of the first polar moiety. This provides addition sites for electrostatic interaction with the second polar moiety of the macrocycle.

In some embodiments, the polymer is a polyamide-imide. In other embodiments, the polymer is polybenzimidazole.

In some embodiments, the macrocycle is sulfocalix[n]arene.

In some embodiments, the macrocycle is 4-sulfocalix[8]arene (SCA8).

In some embodiments, a weight ratio of macrocycle to polymer is about 0.1 wt% to about 5 wt%.

In some embodiments, the membrane composite has a thickness of about 80 pm to about 150 pm.

In some embodiments, the membrane composite has a FTIR spectrum peak of about 1040 cm ¹.

In some embodiments, the membrane composite has a sulphur surface composition of at least about 1 At%.

In some embodiments, the membrane composite has a free volume intensity S- para meter of at least about 0.455.

In some embodiments, the membrane composite has a pore size distribution R- para meter of at least about 0.465.

In some embodiments, the membrane composite has a Victoria blue B rejection of at least about 90% when the macrocycle loading is about 0.1 wt%. In some embodiments, the membrane composite has a methylene blue rejection of at least about 80% when the macrocycle loading is about 0.1 wt%. In some embodiments, the membrane composite has a paracetamol rejection of at least about 30% when the macrocycle loading is about 0.1 wt%.

In some embodiments, the membrane composite has a pure methanol permeance of about 1 LMH/bar to about 1.8 LMH/bar when the macrocycle loading is about 0.3 wt%.

In some embodiments, the membrane composite has a pure acetonitrile permeance of about 1 LMH/bar to about 1.28 LMH/bar when the macrocycle loading is about 0.3 wt%.

In some embodiments, the membrane composite has a pure acetone permeance of about 0.4 LMH/bar to about 1 LMH/bar when the macrocycle loading is about 0.3 wt%.

In some embodiments, the membrane composite has a pure ethanol permeance of about 0.48 LMH/bar to about 0.76 LMH/bar when the macrocycle loading is about 0.3 wt% .

In some embodiments, the membrane composite has a pure ethyl acetate permeance of about 0.24 LMH/bar to about 0.36 LMH/bar when the macrocycle loading is about 0.3 wt% . In some embodiments, the membrane composite has a pure tetrahydrofuran permeance of about 0.12 LMH/bar to about 0.24 LMH/bar when the macrocycle loading is about 0.3 wt% .

In some embodiments, the membrane composite has a pure toluene permeance of about 0.01 LMH/bar to about 0.08 LMH/bar when the macrocycle loading is about 0.3 wt% .

In some embodiments, the membrane composite has a stability in a methylene blue/methanol mixture for at least 7 days. In some embodiments, the membrane composite has a A/,/V-dimethyl-4-nitroaniline (DMNA)/methylene blue (MB) separation factor of about 14.5.

In some embodiments, the membrane composite is formed as a hollow fiber membrane.

The present invention also provides a method of fabricating a membrane composite, the membrane composite having a macrocycle homogenously distributed within a layer of polymer, the method comprising incubating the layer of polymer in a macrocycle solution; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from acyl, amino, acylamino, aminoacyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 5 to 8.

In some embodiments, the method further comprises a step of crosslinking the layer of polymer with 1,6-hexanediamine (HDA), tris(2-aminoethyl)amine (TAEA), a,a'- dibromo-p-xylene or a combination thereof.

In some embodiments, the crosslinking is performed by incubating the layer of polymer in a crosslinker solution.

In some embodiments, the crosslinker solution has a crosslinker concentration of about 5 wt%.

In some embodiments, the crosslinking step is performed for at least 12 h.

In some embodiments, the macrocycle is homogenous dispersed across the thickness of the layer of polymer.

In some embodiments, the macrocycle solution comprises a protic solvent. In some embodiments, the protic solvent is selected from methanol, ethanol, water or a combination thereof.

In some embodiments, the layer of polymer is incubated in a macrocycle solution for at least 0.5 h.

In some embodiments, the method further comprises a drying step after the incubation step.

The present invention also provide a membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; and wherein the macrocycle is characterised by a pore size of more than about 4 A.

Brief description of the drawings

Embodiments of the present invention will now be described, by way of non-limiting example, with reference to the drawings in which:

Figure 1 shows schematics of a conventional method compared to an embodiment of the present invention;

Figure 2 shows confirmation of the molecularly homogeneous infiltration of SCA4 molecules;

Figure 3 shows identification of polymer-SCA4 interactions and the bridging effect on polymer chains;

Figure 4 shows gas transport parameters and sorption behaviors of AOPIM1-SCA4 membranes;

Figure 5 shows illustration of the molecular gatekeeping mechanism in the AOPIM1 microporous scaffold;

Figure 6 shows pure-gas separation performances of AOPIM1-SCA4 membranes compared with recent upper bounds and other high-performance polymer-based membranes; Figure 7 shows molecular structures of a) STCAss and b) SCA, and an exemplary reaction scheme of interfacial polymerization between MPD (aqueous phase) and TMC (oil phase);

Figure 8 shows surface morphologies via FESEM and AFM images of the polyamide layers for a TFC-0, b TFN-STCAss-1.5, and c TFN-SCA-1.5;

Figure 9 shows surface chemistry via ATR-FTIR spectra for TFC-0, TFN-STCAss-1.5, and TFN-SCA-1.5;

Figure 10 shows ionic interactions between the polyamide network and STCAss/SCA. a) Deconvolution of N 1 s and S 2 p XPS spectra for TFC-0, TFN-STCAss-1.5, and TFN- SCA-1.5, and b) illustrations on the interactions between polyamide network and STCAss/SCA;

Figure 11 shows membrane microstructures via a, b) S-parameters and c, d) R- parameters of TFC-0, TFN-STCAss-1.5, and TFN-SCA-1.5. The dotted lines in b indicate the average values of S-parameters of polyamide networks for guiding the view;

Figure 12 shows separation performance and transport properties of TFC-0, TFN- STCAss-1.5, and TFN-SCA-1.5. OSFO performance for a FO mode and b PRO mode, as well as c transport properties and d diffusivities of water, methanol, and ethanol in TFC- 0, TFN-STCAss-1.5, and TFN-SCA-1.5;

Figure 13 shows rejections towards EG, DEG and glucose of the TFC-0, TFN-STCAssl.5, TFN-SCA-1.5;

Figure 14 shows FESEM images of pristine TFC and TFNs with different STCAss loadings; Figure 15 shows optimization of STCAss loading. Ethanol flux (closed square), reverse solute flux (open circle), and Js/Jw (closed triangle) with increasing STCAss loading for a FO mode and b PRO mode. The feed solution is pure ethanol and the draw solution is 2M LiCI in ethanol;

Figure 16 shows surface and cross-sectional morphologies of the as-spun FIFM;

Figure 17 shows membrane morphologies of an exemplary membrane composite (TAEA- SCA8-0.3);

Figure 18 shows ATR-FTIR spectra of the as-spun, FIDA-XIinked, TAEA-XIinked, and TAEA-SCA8-0.3 HFMs;

Figure 19 shows deconvolution of N Is XPS spectra, (a) the as-spun, (b) HDA-XIinked, (c) TAEA-XIinked, and (d) TAEA-SCA8-0.3 HFMs;

Figure 20 is a schematic illustration. The chemical interaction between TAEA-XIinked HFM and SCA8;

Figure 21 shows membrane microstructure, a) S- and b) R-parameters of the as-spun, HDA-XI inked, TAEA-XIinked, and TAEA-SCA8-0.3 HFMs;

Figure 22 shows separation performance of the as-spun, FIDA- and TAEA-XIinked HFMs. a) PWP and MWCO, inset: the molecular structures and sizes of HDA and TAEA, b) pore size distribution and c) pure methanol permeance and rejection towards RB;

Figure 23 shows OSN separation performance. a) Performance of HFMs to separate paracetamol, MB and VBB in methanol as a function of SCA8 loadings, inset: the molecular structures and sizes of paracetamol, MB and VBB. b) Permeances of selected organic solvents through TAEA-SCA-0.3. c) Correlation between organic solvent permeance and solvent properties d) Membrane stability in MB/methanol mixtures for 144 h;

Figure 24 shows a) photographs and UV absorption spectra of the mixed-solute solutions before and after filtration through TAEA-SCA8-0.3, the UV spectra were measured by diluting all the solutions by 3 times b) Illustrations of the molecular sieving mechanism through SCA8 impregnated membrane;

Figure 25A shows the morphology of crosslinked PBI hollow fiber membrane;

Figure 25B shows morphology of crosslinked PBI impregnated with macrocycles hollow fiber membrane;

Figure 26 shows the effect of oil concentration on the permeance and oil rejection (@10 bar 8i room temperature);

Figure 27 shows separation performance of 1-inch XPBI HFM module over time using a cross-flow set-up (10 wt% oil in acetone, at 10 bar & room temperature over 16 days); Figure 28 shows the pore size distribution of PBI hollow fiber membrane;

Figure 29A shows pure solvent permeance of crosslinked PBI-SCA8 hollow fiber membrane at 10 bar and 20 °C; and

Figure 29B shows pure solvent of crosslinked PBI-SCA8 hollow fiber membrane as a function of the inverse of viscosity at 10 bar and 20 °C.

Detailed description

The present invention is predicated on the understanding that, if a nanoporous agent could be incorporated into polymer membranes on a molecularly homogeneous basis after the polymers are fabricated into films or selective layers, the need of undergoing any potentially problematic prefabrication mixing stages could be eliminated (i.e. the direct insertion of entire fillers rather than via the pre-seeding process or in situ syntheses), and at least some of the above mentioned issues can be avoided or at least ameliorated.

Further advantageously, when attempting to solve overcome these problems, the inventors have found that complete solvation of the nanoporous agents in certain solvents can be beneficial, so as to be able to diffuse deep into the fabricated polymer films. The inventors have found that inorganic fillers generally lack good solubility and the organic ones tend to be soluble in solvents that dissolve the polymers too. To this end, both inorganic and organic fillers are disadvantageous. The inventors proposed three key criteria for the choice of material that could possibly enable this design: (1) being water-soluble such that it can be molecularly solvated by common membrane-treating protic solvents, like methanol or ethanol, which can diffuse through the entire polymer microporous structure without dissolving the polymers; (2) being capable of forming extensive electrostatic interactions (such as hydrogen and ionic bonds) with the polymer which are stronger than both the solvation forces and the self-aggregation tendencies in order to be homogeneously distributed and firmly anchored into the membranes (possible to be coincided with criterion (1) because water soluble functionalities tend to form ionic or hydrogen-bond interactions with the N- or O-containing groups not uncommon to find in many polymers); (3) having intrinsic size sieving pores or open cavities to act as molecular-sieving windows for discriminatory gas passage, which will require such materials to have multi-dimensional architectures, but at the same time also possessing a small molecular size for effectively infiltrating the fabricated films without encountering too much mass transfer resistance. Coexistence of these seemingly contradictory criteria had hardly been considered before because the conventionally perceived multi-dimensional nanoporous materials like the various framework-type particles generally lack protic solvent solubility which is, however, more commonly found in small, non-extended molecules. Yet, these small molecules tend not to possess intrinsic pores because they are often 0- or 1- dimensionally structured.

The inventors have found that macrocycles can be used in this regard. For example, the most simply structured member, 4-sulfocalix[4]arene (SCA4), from the organic macrocyclic calixarene family can fit these criteria. SCA4 molecules possess an intrinsic 3-dimensional bowl-shaped cavity (Figure 1) with a small range of bottom opening sizes around a mean value of about 3.0 A resulting from the partially flexible methylene linkers that give rise to conformational flexibility around the bowl shape. Therefore, they can act as size-sieving molecular gatekeepers in membranes that selectively pass gases with a size smaller than or similar to that of its mean bottom opening, like H2 (2.89 A) and CO2 (3.30 A), while strongly impeding larger gas molecules, like N2 (3.64 A) and CFU (3.80 A). Meanwhile, the multiple water-soluble sulfonic groups on the upper rim of SCA4's molecule-sized body enable its complete solvation in methanol so that the solution can carry SCA4 molecules to molecularly infiltrate the microporous structure of already formed polymer films which serves as a microporous scaffold providing active lodging sites with extensive hydrogen and ionic bonding capability. As such, an ultra- facile but unconventional composite membrane design by incorporating porous agents only after the membrane was fabricated was enabled, named herein the post-fabrication infiltration (PFI) membranes. More importantly, besides SCA4, calixarenes are actually a huge class of tuneable and expandable organic porous molecules with other base cavity sizes built with different numbers of p-phenol subunits, like calix[n]arenes, where n can be, but is not limited to, 4, 5, 6 and 8. They are also bestowed with dual sites for functionalization. The upper rim can come with a variety of water-soluble moieties besides sulfonic groups, such as phosphoric or amine groups, that further broaden the pool of PFI-viable polymers, while the lower rim phenolic hydroxyl could also be subjected to chemical modification to provide additional functionalities or finer tuning of the bottom opening size as the cavity can be expanded by appending lower rim substituents. Therefore, this PFI strategy for utilizing molecules of potentially tuneable intrinsic nanoporosity not only reveals a material advancement by delivering high- performance molecular-sieving composite membranes, but also methodological progress by completely bypassing the ingrained issues of interfacial nanodefects and pore blockage in composite membrane designs.

"Alkyl" refers to monovalent alkyl groups which may be straight chained or branched and preferably have from 1 to 10 carbon atoms or more preferably 1 to 6 carbon atoms. Examples of such alkyl groups include methyl, ethyl, n-propyl, iso- propyl, n-butyl, iso- butyl, n-hexyl, and the like.

"Alkenyl" refers to a monovalent alkenyl group which may be straight chained or branched and preferably have from 2 to 10 carbon atoms and more preferably 2 to 6 carbon atoms and have at least 1 and preferably from 1-2, carbon to carbon, double bonds. Examples include ethenyl (-CH =CH2), n-propenyl (-CF-I2CF-I = CH2), /so-propenyl (-C(CH₃)=CH₂), but-2-enyl (-CH₂CH = CHCH₃), and the like.

"Alkoxy" refers to the group alkyl-O- where the alkyl group is as described above. Examples include, methoxy, ethoxy, n-propoxy, /so-propoxy, n-butoxy, tert- butoxy, sec-butoxy, n-pentoxy, n-hexoxy, 1,2-dimethylbutoxy, and the like.

"Alkenyloxy" refers to the group alkenyl-O- wherein the alkenyl group is as described above.

"Oxo/hydroxy" refers to groups =0, HO-.

"Acyl" refers to groups H-C(O)-, alkyl-C(O)-, cycloalkyl-C(O)-, aryl-C(O)-, heteroaryl- C(O)- and heterocyclyl-C(O)-, where alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl are as described herein.

"Oxyacyl" refers to groups HOC(O)-, halide-OC(O)-, alkyl-OC(O)-, cycloalkyl-OC(O)-, aryl-OC(O)-, heteroaryl-OC(O)-, and heterocyclyl-OC(O)-, where alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl are as described herein.

"Acyloxy" refers to the groups -0C(0)-alkyl, -0C(0)-aryl, -C(0)0-heteroaryl, and -C(0)0-heterocyclyl where alkyl, aryl, heteroaryl and heterocyclyl are as described herein.

"Amino" refers to the group -NR"R" where each R" is independently hydrogen, alkyl, cycloalkyl, aryl, heteroaryl, and heterocyclyl and where each of alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl is as described herein.

"Aminoacyl" refers to the group -C(0)NR"R" where each R" is independently hydrogen, alkyl, cycloalkyl, aryl, heteroaryl, and heterocyclyl and where each of alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl is as described herein.

"Acylamino" refers to the group -NR"C(0)R" where each R" is independently hydrogen, alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl and where each of alkyl, cycloalkyl, aryl, heteroaryl, and heterocyclyl are as described herein. "Sulfonyl" refers to groups H-S(0)2-, alkyl-S(0)2-, cycloalkyl-S(0)2-, aryl-S(0)2-, heteroaryl-S(0)2-, and heterocyclyl-S(0)2-, where alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl are as described herein. "Oximyl" refers to groups -CH = N-OH, -C(alkyl) = N-OH, -C(cycloalkyl) = N-OH, - C(aryl) = N-OH, -C(heteroaryl) = N-OH, and -C(heterocyclyl) = N-OH, where alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl are as described herein.

"Amidoximyl" refers to groups -CH = N-OH, -C(alkyl) = N-OH, -C(cycloalkyl) = N-OH, - C(aryl) = N-OH, -C(heteroaryl) = N-OH, and -C(heterocyclyl) = N-OH, where alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl are as described herein.

"Hydrazonyl" refers to groups -CH = N-NH2, -C(alkyl) = N-NH2, -C(cycloalkyl) = N-NH2, - C(aryl) = N-NH2, -C(heteroaryl) = N-NH2, and -C(heterocyclyl) = N-NH2, where alkyl, cycloalkyl, aryl, heteroaryl and heterocyclyl are as described herein.

"Iminyl" refers to groups -CH = N-R, -C(alkyl) = N-R, -C(cycloalkyl) = N-R, -C(aryl) = N-R, -C(heteroaryl) = N-R, and -C(heterocyclyl) = N-R, wherein R is alkyl, cycloalkyl, aryl, heteroaryl or heterocyclyl are as described herein.

In this specification "optionally substituted" is taken to mean that a group may or may not be further substituted or fused (so as to form a condensed polycyclic group) with one or more groups selected from hydroxyl, acyl, alkyl, alkoxy, alkenyl, alkenyloxy, alkynyl, alkynyloxy, amino, aminoacyl, thio, arylalkyl, arylalkoxy, aryl, aryloxy, carboxyl, acylamino, cyano, halogen, nitro, phosphono, sulfo, phosphorylamino, phosphinyl, heteroaryl, heteroarylalkyl, heteroaryloxy, heterocyclyl, heterocyclylalkyl, heterocyclyloxy, oxyacyl, oxime, oxime ether, hydrazone, oxyacylamino, oxysulfonylamino, aminoacyloxy, trihalomethyl, trialkylsilyl, pentafluoroethyl, trifluoromethoxy, difluoromethoxy, trifluoromethanethio, trifluoroethenyl, mono- and di-alkylamino, mono-and di-(substituted alkyl)amino, mono- and di-arylamino, mono- and di-heteroarylamino, mono- and di-heterocyclyl amino, and unsymmetric di- substituted amines having different substituents selected from alkyl, aryl, heteroaryl and heterocyclyl, and the like, and may also include a bond to a solid support material, (for example, substituted onto a polymer resin). For instance, an "optionally substituted amino" group may include amino acid and peptide residues. In chemistry, polarity is a separation of electric charge leading to a molecule or its moieties having an electric dipole moment, with a negatively charged end and a positively charged end. Accordingly, "polar" molecules and "polar" moieties are resultant of a difference in electronegativity between the bonded atoms. Polar molecules and moieties can interact with other molecules and moieties through dipole-dipole intermolecular forces, hydrogen bonds and ionic interactions; i.e. through electrostatic interactions. A hydrogen bond (H-bond) is a primarily electrostatic force of attraction between a hydrogen (H) atom which is covalently bound to a more electronegative atom or group, particularly the second-row elements nitrogen (N), oxygen (0), or fluorine (F)— the hydrogen bond donor (Dn)— and another electronegative atom bearing a lone pair of electrons— the hydrogen bond acceptor (Ac). Such an interacting system is generally denoted Dn-H—Ac, where the solid line denotes a polar covalent bond, and the dotted or dashed line indicates the hydrogen bond. Hydrogen bonding is currently considered as having both covalent and electrostatic contributions.

Ionic bonding is a type of chemical bonding that involves the electrostatic attraction between oppositely charged ions, or between two atoms with sharply different electronegativities. In the simplest case, an ionic bond results from the transfer of electrons from a metal to a non-metal in order to obtain a full valence shell for both atoms. It is important to recognize that clean ionic bonding — in which one atom or molecule completely transfers an electron to another — cannot exist: all ionic compounds have some degree of covalent bonding, or electron sharing. Thus, the term "ionic bonding" is given when the ionic character is greater than the covalent character - that is, a bond in which a large electronegativity difference exists between the two atoms, causing the bonding to be more polar (ionic) than in covalent bonding where electrons are shared more equally. Bonds with partially ionic and partially covalent character are called polar covalent bonds, and are also included within this scope.

An electron withdrawing group refers to an atom or functional group that withdraws electron density from its neighbouring atom(s) or from a conjugated system. Conversely, an electron donating group refers to an atom or functional group that donates electron density to its neighbouring atom(s) or to a conjugated system. A "derivative" is a compound that is derived from a similar compound by a chemical reaction. Derivative also meant a compound that can be imagined to arise from another compound, if one atom or group of atoms is replaced with another atom or group of atoms. Accordingly, calixarene or a derivative thereof can refer to calixarene which has various pendant groups attached to it at various locations of the macrocycle, and can also refer to modifications within the backbone of the macrocycle, such as the number and type of atoms in the linkers connecting the phenol units. Accordingly, the present inventions provides a membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; and wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer.

The membrane composite of the present invention are for use in filtration applications. A membrane is a selective barrier; it allows some things to pass through but stops others. Such things may be, but is not limited to, molecules, ions, gases, solvents, or other small particles.

In some embodiments, the first polar moiety forms an electrostatic interaction with an electron withdrawing group (second polar moiety). In other embodiments, the first polar moiety forms an electrostatic interaction with an electron donating group (second polar moiety). In other embodiments, the first polar moiety is an electron withdrawing group and forms an electrostatic interaction with the second polar moiety. In other embodiments, the first polar moiety is an electron donating group and forms an electrostatic interaction with the second polar moiety. In some embodiments, the electrostatic interaction is selected from dipole-dipole intermolecular force, hydrogen bond, ionic interaction, and a combination thereof. In other embodiments, the electrostatic interaction is selected from hydrogen bond, ionic interaction, and a combination thereof. In some embodiments, the first polar moiety resides in a polymer backbone, as a side group, or a combination thereof. For example, the polar moiety can be an amide linkage. In other embodiments, the polar moiety is resides in a crosslinker residue. For example, the polar moiety can be provided by an amine crosslinker.

In some embodiments, the polar moiety is selected from cyano, acyl, acyloxy, oxyacyl, amino, acylamino, aminoacyl, amidoximyl, oximyl, hydrazonyl, iminyl, hydroxyl, pyrrolyl, oxazolyl, imidazolyl, pyrazolyl, triazolyl, and a combination thereof. In other embodiments, the polar moiety is selected from acyl, amino, acylamino, aminoacyl, amidoximyl, oximyl, iminyl, hydroxyl, and a combination thereof. In other embodiments, the polar moiety is selected from acyl, amino, acylamino, aminoacyl, amidoximyl, oximyl, hydroxyl, and a combination thereof. In other embodiments, the polar moiety is amino and/or acylamino.

In some embodiments, the macrocycle or derivative thereof has an upper rim and a lower rim, the upper rim and lower rim are separated by a frustum-shaped cavity. In other embodiments, a diameter of the lower rim is at least about 3 A, or at least about 4 A, 5 A, 6 A, 7 A, 8 A, or 9 A.

In some embodiments, the upper rim is functionalised with at least one moiety independently selected from sulfonyl, phosphoryl, amino, carboxyl, oxyalkyl and a combination thereof. These moiety can be the second polar moiety. In other embodiments, the upper rim is functionalised with at least one moiety independently selected from sulfonyl, amino, and a combination thereof. Advantageously, this provides the polar moiety in order to form an electrostatic interaction with the polar moiety of the polymer. These moieties can further be optionally substituted. For example, the sulfonyl moiety can be optionally substituted with alkyl groups, which does not affect (or minimally affect) the polarity of the sulfonyl group.

In some embodiments, the upper rim is functionalised with at least 2 moieties. In other embodiments, the upper rim is functionalised with 3, 4, 5 or 6 moieties.

The moieties on the upper rim and/or lower rim of the macrocycle or derivative thereof can form the second polar moiety. Alternatively, as mentioned above, the moieties on the upper rim and/or lower rim of the macrocycle or derivative thereof can improve the macrocyle's water-solublility such that it can be molecularly solvated by common membrane-treating protic solvents and diffuse through the entire polymer microporous structure without dissolving the polymers.

The first polar moiety and the second polar moiety are for forming an electrostatic interaction between them. Thus, for example, if the first polar moiety has a positive polarity, the second polar moiety must have a negative polarity.

In some embodiments, the lower rim is functionalised with at least one moiety independently selected from optionally substituted acyloxy, optionally substituted oxyacyl, optionally substituted alkyl, optionally substituted alkenyl, optionally substituted alkoxy, and optionally substituted alkenyloxy. In other embodiments, the lower rim is functionalised with at least one moiety independently selected from C1-C5 alkyl, C2-C5 alkenyl, C1-C5 alkoxy, and C2-C5 alkenyloxy. Advantageously, by varying the chain length of the moiety, the pore size of the macrocycle can be varied.

For example, when the first polar moiety is amidoximyl and/or amidyl, the second polar moiety can be a sulfo moiety.

In some embodiments, the macrocycle is selected from cyclodextrin, calixarene, cucurbituril, resorcinarene, pillararene, and a combination thereof.

In some embodiments, the macrocycle or derivative thereof is a calixarene or a derivative thereof. In other embodiments, the calixarene is selected from sulfocalixarene (or sulfonylcalixarene), sulfothiacalixarene, carboxylatocalixarene, aminocalixrene, p-phosphonic acid calixarene, or a combination thereof.

In some embodiments, the macrocycle or derivative thereof has monomer residues of 4 to 12. In other embodiments, the number of monomer residues is 4, 5, 6, 7, 8, 9, 10, 11 or 12.

In some embodiments, a mole ratio of macrocycle to polymer is about 0.1% to about 10%. In other embodiments, the mole ratio is about 0.1% to about 9%, about 0.1% to about 8%, about 0.1% to about 7%, about 0.1% to about 6%, about 0.1% to about 5%, about 0.1% to about 4%, or about 0.1% to about 3%. In some embodiments, the macrocycle is about 0.1 wt% to about 10 wt% relative to the membrane composite. In other embodiments, the relative weight is about 0.1% to about 9%, about 0.1% to about 8%, about 0.1% to about 7%, about 0.1% to about 6%, about 0.1% to about 5%, about 0.1% to about 4%, or about 0.1% to about 3%.

In some embodiments, the macrocycle is homogenously dispersed across the thickness of the layer of polymer. Accordingly, when the macrocycles are allowed to infiltrate the polymer layer, the macrocycles penetrates within the polymer layer and is evenly distributed within the polymer layer. In some embodiments, the macrocycles are also not aggregated or not agglomerated.

The polymer in the layer of polymer can be a polymer with intrinsic microporosity-1 (PIM-1), polyamide, polyamide-imide, polybenzimidazole, polylactic acid (PLA), polycarboxyl ic acid (PCA), polyethylenimine (PEI), polyarylamine, polyalkylamine, polyallylamine, poly(vinyl amine) or a combination thereof.

The polymer in the layer of polymer can be further crosslinked. The present invention also provides a method of fabricating a membrane composite, the membrane composite having a macrocycle homogenously distributed within a layer of polymer, the method comprising incubating the layer of polymer in a macrocycle solution; wherein the polymer comprises a first polar moiety; and wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer.

In some embodiments, the electrostatic interaction is selected from hydrogen-bond, proton transfer, ionic interaction or a combination thereof. Advantageously, by using electrostatic interaction, post fabrication infiltration of macrocycles into a polymer can be utilised which can better streamline membrane fabrication processes. As no additional surfactants or linkers are used, the cost of fabrication is also reduced. In other embodiments, the macrocycle is further covalently bonded to the polymer. This covalent bonding of the macrocycle to the polymer can be performed either as a step after the post fabrication infiltration or during the post fabrication infiltration, in which the macrocycle is electrostatically interacted with the polymer. Advantageously, this further improves the stability of the membrane composite.

In some embodiments, the macrocycle solution comprises a protic solvent. A protic solvent is a solvent that has a hydrogen atom bound to an oxygen (as in a hydroxyl group), a nitrogen (as in an amine group), or fluoride (as in hydrogen fluoride). In general terms, any solvent that contains a labile H⁺ is called a protic solvent. The molecules of such solvents readily donate protons (H⁺) to solutes, often via hydrogen bonding. Water is the most common protic solvent. Conversely, polar aprotic solvents cannot donate protons but still have the ability to dissolve many salts.

In some embodiments, the protic solvent is selected from methanol, ethanol, isopropanol, n-propanol, water, and a combination thereof. In other embodiments, the protic solvent is selected from butanol, isopropanol, nitromethane, acetic acid, formic acid and a combination thereof.

The present invention is herein exemplified in three applications. In a first embodiment, the membrane composite comprises amidoxime-functionalised intrinsic microscopic polymer (AOPIM-1) and 4-sulfocalix[4]arene (SCA4). The polymer is synthesised from 5,5',6,6'-tetrahydroxy-3,3,3,3'-tetramethyl-l,l'-spirobisindane (TTSBI) and 2, 3,5,6- tetra-fluoroterephthalonitrile (TFTPN). This membrane composite can be used in the separation of gases. In a second embodiment, the membrane composite comprises a polyamide and sulfothiacalix[4]arene (STCAss) or sulfocalix[4]arene (SCA). The polyamide can be formed using interfacial polymerisation in the presence of macrocycles. This membrane composite can be further casted on a substrate, such as Matrimid. This membrane composite can be used in organic solvent forward osmosis. In a third embodiment, the membrane composite comprises polyamide-imide (Torlon) and 4-sulfocalix[8]arene (SCA8). The membrane composite can be further crosslinked using amine crosslinkers. The membrane composite can be formed as a hollow fiber membrane and used in organic solvent nanofiltration. In a further example, the membrane composite comprises polybenzimidazole and 4- sulfocalix[8]arene. The polymer can be crosslinked using xylene crosslinkers. Accordingly, the membrane composite can be used as a filtration membrane. In some embodiments, the membrane composite is for use in gas separation. The gas can be a mixture of Hi, Oi, N2, CFU, and CO2. In other embodiments, the membrane composite is for use in organic solvent filtration. For example, in organic solvent forward osmosis, the organic solvent can flow from less concentrated solution across the membrane to the more concentrated solution. This allows for the concentration of salts, molecules or other non-permeable solvents. In other embodiments, the membrane composite is for use in organic solvent nanofiltration, in which organic solvents are allowed to pass through while solutes in the range of 200 g/mol to 1000 g/mol are retained. In some embodiments, the permeance is based on the molecular size of the solvent or gas. Towards this end, the size selectivity of the membrane composite can be tuned by altering the type (and hence cavity size) of the macrocycle. In other embodiments, the permeance is based on the viscosity of the solvent. A low viscosity tends to have a high permeance across the membrane composite. In other embodiments, the permeance is based on the affinity of the solvent or gas towards the membrane composite. The affinity can be tuned by varying the type of polymer in the polymer layer and macrocycle, in consideration of the polarity and/or protic nature of the solvent and gas. A high affinity towards the membrane composite tends to exhibit a high permeance across the membrane. In other embodiments, the permeance is linearly correlated to σμ^-1V^-1 (solubility parameter (S), viscosity (μ) and molar volume (V) of the solvent).

Accordingly, the present invention also provides a membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from cyano, amino, amidoximyl, oximyl, hydroxyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 4 to 8.

In some embodiments, the membrane composite comprises: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises an electron withdrawing moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from cyano, amino, amidoximyl, oximyl, hydroxyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 4 to 8.

In some embodiments, the polymer comprises polymerised monomer residues of Formula (I):

wherein Ri and R2 are independently selected from acylamino, aminoacyl, oxyacyl, acyloxy, cyano, amino, amidoximyl, oximyl, hydroxy; and wherein m is an integer between 40 and 450.

In some embodiments, Ri and R2 are independently selected from amino, amidoximyl, oximyl, and hydroxyl. In other embodiments, Ri and Rå are independently selected from amino, amidoximyl, and oximyl. In other embodiments, Ri and R2 are both amidoximyl. The value m refers to the number of polymerised monomeric units in the polymer. In some embodiments, m is an integer between 60 and 450, 80 and 450, 80 and 400, 80 and 350, 120 and 350, 120 and 300, 120 and 250, 120 and 200, or 150 and 200. The value m can be obtained by gel permeation chromatography, based on weight-average molecular weight (M_w). In general, the polymerisation results in a Gaussian distribution of molecular weights. Accordingly, the skilled person would understand that a range of m values is to be expected. In particular, it was found to be further advantageous when m is larger than 40.

The degree of polymerization, or DP, is the number of monomeric units in a macromolecule or polymer or oligomer molecule, and can be calculated by dividing M_n (number-average molecular weight) over Mo (molecular weight of the monomer unit). In some embodiments, DP is about 20 to about 250. In other embodiments, DP is about 40 to about 250, about 60 to about 250, about 60 to about 230, about 60 to about 210, about 80 to about 210, about 80 to about 190, about 80 to about 170, about 80 to about 150, or about 100 to about 150.

In some embodiments, the polymer is an intrinsically microporous polymer. Polymers of intrinsic microporosity (PIM) are a unique class of microporous material. PIMs contain a continuous network of interconnected intermolecular voids less than 2 nm in width. Classified as a porous organic polymer, PIMs generate porosity from their rigid and contorted macromolecular chains that do not efficiently pack in the solid state. PIMs are composed of a fused ring sequences interrupted by Spiro-centers or other sites of contortion along the backbone. Due to their fused ring structure PIMs cannot rotate freely along the polymer backbone, ensuring the macromolecular components conformation cannot rearrange and ensuring the highly contorted shape is fixed during synthesis. Some examples of functionalized PIM-1 are AO-PIM-1, TZPIM, Thio-PIM-1, PIM-l-COOH, which are shown below.

In some embodiments, the macrocycle is sulfocalix[n]arene or a derivative thereof. In some embodiments, the macrocycle is sulfocalix[4]arene. In some embodiments, the calix[n]arene or a derivative thereof has an upper rim and a lower rim, the upper rim and lower rim are separated by a frustum-shaped cavity, wherein a diameter of the lower rim is at least about 3 A.

In some embodiments, the upper rim is functionalised with at least one water-soluble moiety independently selected from sulfonyl, phosphoryl, amino, and a combination thereof. As used herein, "water-soluble moiety" refers to a moiety the presence of which causes the molecule to be water soluble. In other embodiments, the at least one water- soluble moiety independently selected from sulfonyl, amino, and a combination thereof. In other embodiments, the at least one water-soluble moiety independently is sulfonyl.

In some embodiments, the upper rim is functionalised with at least 4 water-soluble moieties. In other embodiments, the upper rim is functionalised with at least 1, 2, 3, 5, 6, 7 or 8 water-soluble moieties. In some embodiments, the lower rim is functionalised with at least one moiety independently selected from alkyl, alkenyl, alkoxy, and alkenyloxy. In other embodiments, the lower rim is functionalised with at least one moiety independently selected from alkyl, and alkoxy. In other embodiments, the lower rim is functionalised with at least one moiety independently selected from C1-C5 alkyl, and C1-C5 alkoxy.

In some embodiments, a mole ratio of macrocycle to polymer is about 1% to about 5%. In other embodiments, the mole ratio is about 0.1% to about 5%, about 0.1% to about 4%, about 1% to about 4%, about 0.1% to about 3%, or about 1% to about 3%. In some embodiments, the macrocycle is homogenously dispersed across the thickness of the layer of polymer.

In some embodiments, the membrane composite has a smooth surface. In some embodiments, the membrane composite has a thickness of about 20 mm to about 50 mm. In other embodiments, the thickness is about 10 mm to about 50 mm, about 10 mm to about 40 mm, or about 20 mm to about 40 mm.

In some embodiments, the membrane composite has a first d-spacing of about 6.2 A to about 6.4 A, and a second d-spacing of about 4.85 A to about 4.9 A.

In some embodiments, the membrane composite has a tensile strength of about 33 MPa to about 35 MPa.

In some embodiments, the membrane composite has a first free volume radii of about 1.9 A to about 2.1 A, and a second free volume radii of about 3.4 A to about 3.6 A.

In some embodiments, the membrane composite has a PI 2 permeability of about 540 Barrer to about 870 Barrer at 35 °C and 3.5 bar. In other embodiments, the PI 2 permeability is about 560 Barrer to about 870 Barrer, about 580 Barrer to about 870 Barrer, about 600 Barrer to about 870 Barrer, about 620 Barrer to about 870 Barrer, about 640 Barrer to about 870 Barrer, about 660 Barrer to about 870 Barrer, about 680 Barrer to about 870 Barrer, about 700 Barrer to about 870 Barrer, about 720 Barrer to about 870 Barrer, about 740 Barrer to about 870 Barrer, about 760 Barrer to about 870 Barrer, about 780 Barrer to about 870 Barrer, or about 800 Barrer to about 870 Barrer.

In some embodiments, the membrane composite has a O2 permeability of about 30 Barrer to about 140 Barrer at 35 °C and 3.5 bar. In other embodiments, the O2 permeability is about 50 Barrer to about 140 Barrer, about 70 Barrer to about 140 Barrer, about 90 Barrer to about 140 Barrer, or about 110 Barrer to about 140 Barrer.

In some embodiments, the membrane composite has having a N2 permeability of about 4 Barrer to about 30 Barrer at 35 °C and 3.5 bar. In other embodiments, the N2 permeability is about 6 Barrer to about 30 Barrer, about 8 Barrer to about 30 Barrer, about 10 Barrer to about 30 Barrer, about 15 Barrer to about 30 Barrer, about 20 Barrer to about 30 Barrer, or about 25 Barrer to about 30 Barrer.

In some embodiments, the membrane composite has a CPU permeability of about 1 Barrer to about 30 Barrer at 35 °C and 3.5 bar. In other embodiments, the CPU permeability is about 5 Barrer to about 30 Barrer, about 10 Barrer to about 30 Barrer, about 15 Barrer to about 30 Barrer, about 20 Barrer to about 30 Barrer, or about 25 Barrer to about 30 Barrer.

In some embodiments, the membrane composite has a CO2 permeability of about 150 Barrer to about 720 Barrer at 35 °C and 3.5 bar. In other embodiments, the CO2 permeability is about 200 Barrer to about 720 Barrer, about 250 Barrer to about 720 Barrer, about 300 Barrer to about 720 Barrer, about 350 Barrer to about 720 Barrer, about 400 Barrer to about 720 Barrer, about 450 Barrer to about 720 Barrer, about 500 Barrer to about 720 Barrer, about 550 Barrer to about 720 Barrer, about 600 Barrer to about 720 Barrer, or about 650 Barrer to about 720 Barrer.

In some embodiments, the membrane composite has a H2/N2 selectivity of about 30 to about 120 at 35 °C and 3.5 bar. In other embodiments, the H2/N2 selectivity is about 50 to about 120, about 70 to about 120, or about 90 to about 120.

In some embodiments, the membrane composite has a H2/CH4 selectivity of about 30 to about 300 at 35 °C and 3.5 bar. In other embodiments, the H2/CH4 selectivity of about 50 to about 300, about 100 to about 300, about 150 to about 300, about 200 to about 300, or about 250 to about 300.

In some embodiments, the membrane composite has a H2/CO2 selectivity of about 1 to about 4 at 35 °C and 3.5 bar. In other embodiments, the H2/CO2 selectivity is about 2 to about 4, or about 3 to about 4.

In some embodiments, the membrane composite has a O2/N2 selectivity of about 5 to about 8 at 35 °C and 3.5 bar. In other embodiments, the O2/N2 selectivity is about 6 to about 8, or about 7 to about 8.

In some embodiments, the membrane composite has a CO2/N2 selectivity of about 20 to about 40 at 35 °C and 3.5 bar. In other embodiments, the CO2/N2 selectivity is about 25 to about 40, about 30 to about 40, or about 35 to about 40.

In some embodiments, the membrane composite has a CO2/CH4 selectivity of about 30 to about 90 at 35 °C and 3.5 bar. In other embodiments, the CO2/CH4 selectivity is about 30 to about 90, about 40 to about 90, about 50 to about 90, about 60 to about 90, about 70 to about 90, or about 80 to about 90.

In some embodiments, the membrane composite has a stability against physical aging for at least 60 days.

The present invention also provides a method of fabricating a membrane composite, the membrane composite having a macrocycle homogenously distributed within a layer of polymer, the method comprising incubating the layer of polymer in a macrocycle solution; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from cyano, amino, amidoximyl, oximyl, hydroxyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 4 to 8.

In some embodiments, the electrostatic interaction is selected from hydrogen-bond, proton transfer interaction or a combination thereof.

In some embodiments, the macrocycle solution comprises a protic solvent.

In some embodiments, the protic solvent is selected from methanol, ethanol, or a combination thereof.

In some embodiments, the layer of polymer is incubated in a macrocycle solution for at least 12 h. In other embodiments, the incubation is for at least 24 h.

In some embodiments, the method further comprises a drying step after the incubation step. The present invention also provides a membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety, wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from amino, acyl, acylamino, aminoacyl, a combination thereof; wherein the macrocycle is a calix[n]arene or a derivative thereof; and wherein n is an integer selected from 4 to 8.

In some embodiments, the membrane composite comprises: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety, wherein the macrocycle comprises an electron withdrawing moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from amino, acyl, acylamino, aminoacyl, a combination thereof; wherein the macrocycle is a calix[n]arene or a derivative thereof; and wherein n is an integer selected from 4 to 8.

In some embodiments, the polymer comprises polymerised monomer residues of: a) at least one amino monomer having at least two amine groups/moieties; and b) at least one oxyacyl monomer having at least two carboxyl halide groups/moieties.

In some embodiments, the at least one amino monomer having at least two amine groups is selected from m-phenylenediamine (MPD), p-phenylenediamine, o- phenylenediamine, benzenetriamine, benzenehexamine, alkyldiamine, alkyltriamine, alkyltetramine, aryldiamine, aryltriamine, and a combination thereof.

In some embodiments, the at least one acyloxy monomer having at least three carboxyl halide groups. In other embodiments, the at least one acyloxy monomer having at least two carboxyl halide groups is trimesoyl chloride (TMC), terephthaloyl chloride, isopthaloyl chloride, optionally substituted diacyl halide, optionally substituted triacyl halide, and a combination thereof. In other embodiments, the diacyl halide and/or triacyl halide is optionally substituted with alkyl and/or aryl. In other embodiments, the diacyl halide and/or triacyl halide is optionally substituted with C1-C5 alkyl, phenyl, or benzyl. For example, the diacyl halide and/or triacyl halide can be trimethyl acetyl chloride, triphenylacetyl chloride or the likes.

In some embodiments, a weight ratio of MPD to TMC is about 10: 1 to about 40:1. In other embodiments, the weight ratio is about 10: 1 to about 35: 1, about 10: 1 to about 30: 1, about 10: 1 to about 25: 1, or about 10: 1 to about 20: 1.

In some embodiments, the polymer is a polyamide.

In some embodiments, the calix[n]arene or a derivative thereof has an upper rim and a lower rim, the upper rim and lower rim are separated by a frustum-shaped cavity, wherein a diameter of the lower rim is at least about 3 A.

In some embodiments, the upper rim is functionalised with at least one moiety independently selected from sulfonyl, phosphoryl and amino. In other embodiments, the upper rim is functionalised with at least one moiety independently selected from sulfonyl, and amino. In other embodiments, the upper rim is functionalised with at least one moiety selected from sulfonyl.

In some embodiments, the upper rim is functionalised with at least 1, 2, 3, 4, 5, 6, 7 or 8 moieties.

In some embodiments, the lower rim is functionalised with at least one moiety independently selected from alkyl, alkenyl, alkoxy, and alkenyloxy. In other embodiments, the lower rim is functionalised with at least one moiety independently selected from alkyl, and alkoxy. In other embodiments, the lower rim is functionalised with at least one moiety independently selected from C1-C5 alkyl, and C1-C5 alkoxy.

In some embodiments, the macrocycle is selected from sulfothiacalix[n]arene, sulfocalix[n]arene or a derivative thereof "n" can be a integer from 4 to 12, preferentially from 4 to 8. In some embodiments, the macrocycle is selected from sulfothiacalix[4]arene and sulfocalix[4]arene.

In some embodiments, a weight ratio of macrocycle relative to the membrane composite is about 0.1 wt% to about 2 wt%. In other embodiments, the weight ratio is about 0.1 wt% to about 1.5 wt%, or about 0.1 wt% to about 1.0 wt%.

In some embodiments, the membrane composite has a Fourier transform infrared spectroscopy (FTIR) peak of about 1149 cm^-1.

In some embodiments, the membrane composite has a sulphur surface composition of at least about 0.2 At%.

In some embodiments, the membrane composite has a free volume intensity S- parameter of about 0.3 to about 0.6. In other embodiments, the S-parameter is about 0.35 to about 0.6, about 0.35 to about 0.55, about 0.4 to about 0.55, or about 0.4 to about 0.5.

In some embodiments, the membrane composite has a pore size distribution R- parameter of about 0.3 to about 0.7. In other embodiments, the R-parameter is about 0.35 to about 0.7, about 0.35 to about 0.65, about 0.4 to about 0.65, about 0.4 to about 0.6, about 0.45 to about 0.6, about 0.45 to about 0.55, or about 0.45 to about 0.5.

In some embodiments, the membrane composite has surface roughness RMS average roughness (R_q) of about 40 nm to about 90 nm.

In some embodiments, the membrane composite has surface roughness average roughness (R_a) of about 30 nm to about 70 nm.

In some embodiments, the membrane composite has an ethanol flux (Jw) of about 1.8 Lrrr² h^_1 (LMH) to about 4 LMFI in a FO mode (the membrane composite facing a feed solution). In other embodiments, the ethanol flux (Jw) is about 2 LMFI to about 4 LMFI, about 2.5 LMFI to about 4 LMFI, about 3 LMFI to about 4 LMFI, or about 4.5 LMFI to about 4 LMH. In some embodiments, the membrane composite has a reverse solute flux (J_s) of about 0.1 grrr² h^_1 (gMH) to about 0.9 gMH in a FO mode (the membrane composite facing a feed solution). In other embodiments, the reverse solute flux (J_s) is about 0.2 gMH to about 0.9 gMH, about 0.3 gMH to about 0.9 gMH, about 0.4 gMH to about 0.9 gMH, about 0.5 gMH to about 0.9 gMH, about 0.6 gMH to about 0.9 gMH, about 0.7 gMH to about 0.9 gMH, or about 0.8 gMH to about 0.9 gMH.

In some embodiments, the membrane composite has a J_s/Jw of about 0.05 to about 0.4 in a FO mode (the membrane composite facing a feed solution). In other embodiments, the Js/Jw is about 0.1 to about 0.4, about 0.2 to about 0.4, or about 0.3 to about 0.4.

In some embodiments, the membrane composite has an ethanol flux (Jw) of about 2.6 LMH to about 6 LMH in a PRO mode (the membrane composite facing a draw solution). In other embodiments, the ethanol flux (Jw) is about 3 LMH to about 6 LMH, about 3.5 LMH to about 6 LMH, about 4 LMH to about 6 LMH, about 4.5 LMH to about 6 LMH, or about 5 LMH to about 6 LMH.

In some embodiments, the membrane composite has a reverse solute flux (J_s) of about 0.4 gMH to about 4 gMH in a PRO mode (the membrane composite facing a draw solution). In other embodiments, the reverse solute flux (J_s) is about 0.8 gMH to about 4 gMH, about 1 gMH to about 4 gMH, about 1.5 gMH to about 4 gMH, about 2 gMH to about 4 gMH, about 2.5 gMH to about 4 gMH, or about 3 gMH to about 4 gMH. In some embodiments, the membrane composite has a J_s/Jw of about 0.05 to about 0.3 in a PRO mode (the membrane composite facing a draw solution). In other embodiments, the Js/Jw is about 0.1 to about 0.3, about 0.15 to about 0.3, about 0.2 to about 0.3, or about 0.25 to about 0.3. In some embodiments, the membrane composite has a pure ethanol permeance about 0.1 LMH/bar to about 0.4 LMH/bar in a dead-end filtration mode. In other embodiments, the pure ethanol permeance is about 0.2 LMH/bar to about 0.4 LMH/bar, or about 0.3 LMH/bar to about 0.4 LMH/bar. In some embodiments, the membrane composite has a salt rejection of at least about 99%. In other embodiments, the salt rejection is at least about 99.2%, about 99.4%, about 99.5%, about 99.6%, about 99.7%, or about 99.8%.

In some embodiments, the membrane composite has a water diffusivity of about 2.6 x 10^'9 cm²/s to about 3 x 10^-9 cm²/s. In other embodiments, the water diffusivity is about 2.7 x 10^-9 cm²/s to about 3 x 10^-9 cm²/s, about 2.8 x 10^-9 cm²/s to about 3 x 10^-9 cm²/s, or about 2.9 x 10^-9 cm²/s to about 3 x 10^-9 cm²/s.

In some embodiments, the membrane composite has a methanol diffusivity of about 1.2 x 10^"9 cm²/s to about 1.5 x 10 ⁹ cm²/s. In other embodiments, the methanol diffusivity is about 1.3 x 10^-9 cm²/s to about 1.5 x 10^-9 cm²/s, or about 1.4 x 10^-9 cm²/s to about 1.5 x 10^'9 cm²/s.

In some embodiments, the membrane composite has an ethanol diffusivity of about 0.1 x 10^-9 cm²/s to about 1 x 10⁹ cm²/s. In other embodiments, the ethanol diffusivity is about 0.2 x 10^-9 cm²/s to about 1 x 10^-9 cm²/s, about 0.4 x 10^-9 cm²/s to about 1 x 10- ⁹ cm²/s, about 0.6 x 10^-9 cm²/s to about 1 x 10^-9 cm²/s, or about 0.8 x 10^-9 cm²/s to about 1 x 10^-9 cm²/s.

In some embodiments, the membrane composite has a rejection of ethylene glycol of at least about 50%. In other embodiments, the rejection of ethylene glycol is at least about 60%, about 70%, about 80%, about 90%, or about 95%.

In some embodiments, the membrane composite has a rejection of diethylene glycol of at least about 90%. In other embodiments, the rejection of diethylene glycol is at least about 92%, about 94%, about 96%, about 98%, or about 99%.

The present invention also provides a method of fabricating a membrane composite, the membrane composite comprising a macrocycle homogenously distributed within a layer of polymer on a substrate, the method comprising : a) incubating the substrate in a macrocycle solution in order to form a film on a surface of the substrate, the macrocycle solution further comprising at least one amino monomer having at least two amine groups; b) depositing a monomer solution comprising at least one oxyacyl monomer having at least two carboxyl halide groups on the film; and c) crosslinking the at least one amino monomer having at least two amine groups and the at least one oxyacyl monomer having at least two carboxyl halide groups in order to form the layer of polymer; wherein the polymer comprises a first polar moiety, wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from amino, acyl, acylamino, aminoacyl, a combination thereof; wherein the macrocycle is a calix[n]arene or a derivative thereof; and wherein n is an integer selected from 4 to 8.

In some embodiments, the macrocycle solution further comprises an aqueous medium.

The term "aqueous solution" used herein refers to a water based solvent or solvent system, and which comprises of mainly water. Such solvents can be either polar or nonpolar, and/or either protic or aprotic. Solvent systems refer to combinations of solvents which resulting in a final single phase. Both 'solvents' and 'solvent systems' can include, and is not limited to, pentane, cyclopentane, hexane, cyclohexane, benzene, toluene, dioxane, chloroform, diethylether, dichloromethane, tetrahydrofuran, ethyl acetate, acetone, dimethylformamide, acetonitrile, dimethyl sulfoxide, nitromethane, propylene carbonate, formic acid, butanol, isopropanol, propanol, ethanol, methanol, acetic acid, ethylene glycol, diethylene glycol or water. Water based solvent or solvent systems can also include dissolved ions, salts and molecules such as amino acids, proteins, sugars and phospholipids. Such salts may be, but not limited to, sodium chloride, potassium chloride, ammonium acetate, magnesium acetate, magnesium chloride, magnesium sulfate, potassium acetate, potassium chloride, sodium acetate, sodium citrate, zinc chloride, HEPES sodium, calcium chloride, ferric nitrate, sodium bicarbonate, potassium phosphate and sodium phosphate. As such, biological fluids, physiological solutions and culture medium also fall within this definition. In most embodiments, the aqueous solution is water. In some embodiments, the aqueous solution is deionised water. In some embodiments, the aqueous solution is Millipore water.

In some embodiments, the monomer solution further comprises an organic medium. In some embodiments, the at least one amino monomer having at least two amine groups is about 2 wt% relative to the macrocycle solution. In other embodiments, the at least one amino monomer having at least two amine groups is about 1 wt% to about 10 wt%, about 1 wt% to about 9 wt%, about 1 wt% to about 8 wt%, about 1 wt% to about 7 wt%, about 1 wt% to about 6 wt%, about 1 wt% to about 5 wt%, about 1 wt% to about 4 wt%, or about 1 wt% to about 3 wt%.

In some embodiments, the macrocycle solution further comprises sodium dodecyl sulfate at about 0.2 wt% relative to the macrocycle solution. In other embodiments, sodium dodecyl sulfate is about 0.1 wt% to about 0.5 wt%, about 0.1 wt% to about 0.4 wt%, or about 0.1 wt% to about 0.3 wt%.

In some embodiments, the substrate is incubated in the macrocycle solution for at least about 2 min. In other embodiments, the substrate is incubated in the macrocycle solution for at least about 3 min, 4 min, 5 min, or 10 min.

In some embodiments, the at least one oxyacyl monomer having at least two carboxyl halide groups is about 0.1 wt% relative to the monomer solution. In other embodiments, the at least one oxyacyl monomer having at least two carboxyl halide groups is about 0.1 wt% to about 1 wt%, about 0.1 wt% to about 0.9 wt%, about 0.1 wt% to about 0.8 wt%, about 0.1 wt% to about 0.7 wt%, about 0.1 wt% to about 0.6 wt%, about 0.1 wt% to about 0.5 wt%, about 0.1 wt% to about 0.4 wt%, or about 0.1 wt% to about 0.3 wt%.

In some embodiments, the monomer solution is deposited on the film for at least about 1 min.

In some embodiments, the at least one oxyacyl monomer having at least two carboxyl halide groups is selected from trimesoyl chloride (TMC), terephthaloyl chloride, isopthaloyl chloride, optionally substituted diacyl halide, optionally substituted triacyl halide, and a combination thereof.

In some embodiments, a weight ratio of the at least one amino monomer having at least two amine groups to the at least one oxyacyl monomer having at least two carboxyl halide groups is about 10: 1 to about 40:1. In other embodiments, the weight ratio is about 10: 1 to about 35: 1, about 10: 1 to about 30:1, about 10: 1 to about 25: 1, about 10: 1 to about 20: 1, or about 10: 1 to about 15: 1. The weight ratio can be a weight ratio of MPD:TMC.

In some embodiments, a mole ratio of the at least one amino monomer having at least two amine groups to the at least one oxyacyl monomer having at least two carboxyl halide groups is about 10: 1 to about 40: 1. In other embodiments, the mole ratio is about 10: 1 to about 35: 1, about 10: 1 to about 30:1, about 10: 1 to about 25: 1, about 10: 1 to about 20: 1, or about 10: 1 to about 15: 1. The mole ratio can be a weight ratio of MPD:TMC.

In some embodiments, the membrane composite comprises: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises an electron withdrawing moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from acyl, amino, acylamino, aminoacyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 5 to 8.

Advantageously, when these macrocycles with n being 5 or larger are used, it was found that the membrane composite are especially beneficial for use in organic solvent filtration. It was found that using a macrocycle with these monomeric unit allows for the tuning of the pore size in the membrane composite. In fact, a pore size of up to 9.2 A can be obtained if n = 8. This is in contrast to earlier teachings, which suggests that small macrocycles should be used to fabricate a membrane composite for gas separation. Another work suggests that the backbone structure of the macrocycle should be altered to accommodate the permeance of solvents. In that work, it was shown that the pore size can only be enlarged up to about 5 A, which may not be sufficient for larger solvent molecules.

In some embodiments, the layer of polymer is crosslinked with a crosslinker having at least 2 pendant groups, the pendant group selected from aminoalkyl, haloalkyl, and a combination thereof. In other embodiments, the crosslinker is an amine crosslinker having at least 2 pendant aminoalkyl groups. In other embodiments, the amine crosslinker is selected from tris(2-amino)ethylamine (TAEA), 1,6-hexanediamine (HDA), or a combination thereof. In other embodiments, the crosslinker is an aryl crosslinker having at least 2 pendant haloalkyl groups. In other embodiments, the aryl crosslinker is a,a'-dibromo-p-xylene. Accordingly, the crosslinker can be selected from tris(2- amino)ethylamine (TAEA), 1,6-hexanediamine (HDA), a,a'-dibromo-p-xylene or a combination thereof. In other embodiments, the crosslinker is selected from a, a' - dichloro-p-xylene (DCX), 1,4-dibromobutane (DBB), a, a' -dibromo-p-xylene (DBX), glutaraldehyde (GA), 2, 7, 8-diepoxyoctane (DEO), terephthaloyl chloride (TCL), and propargyl bromide.

In some embodiments, the crosslinker comprises a polar moiety. When incorporated into and crosslinked with the polymer, the polar moiety of the crosslinker can form a portion of the first polar moiety. In other embodiments, the polar moiety of the crosslinker forms all of the first polar moiety of the polymer layer. In this regard, the polymer before it is crosslinked does not contain a polar moiety.

Advantageously, this provides addition sites for electrostatic interaction with the second polar moiety of the macrocycle. Further advantageously, use of such crosslinkers also provides a platform for polymer without polar moieties to be used.

In some embodiments, the polymer is a polyamide-imide. Polyamide-imides are either thermosetting or thermoplastic, amorphous polymers that have exceptional mechanical, thermal and chemical resistant properties. They can be prepared from isocyanates and TMA (trimellic acid-anhydride) in N-methyl-2-pyrrolidone (NMP). A prominent distributor of polyamide-imides is Solvay Specialty Polymers, which uses the trademark Torlon. In some embodiments, the polymer is polybenzimidazole (PBI).

In some embodiments, the macrocycle is sulfocalix[n]arene or a derivative thereof. In some embodiments, n is an integer selected from 5 to 8. In some embodiments, the macrocycle is 4-sulfocalix[8]arene (SCA8).

It was found that SCA8 is particularly advantageous over other calixarene such as SCA4 or STCAss. Without wanting to be bound by theory, it is believed that the size of SCA8 is similar to the pore size of the polymer layer, such that when the ionic bonds are formed, the macrocycle structure is not strained or warped. This allows for a consistent permeance and rejection when used as a filtration membrane. Further, SCA8 also provides a larger cavity size for the permeance of solvents. While a larger cavity size will also allow some solutes to pass through, this was advantageously not found to be case as the polymer layer also provides a tortuous path which hinders the permeance of the solutes.

For example, polybenzimidazole (PBI) can be used as the polymeric material to synthesize integrally skinned asymmetric OSN FIFMs in view of its excellent mechanical and chemical characteristics, which are attributed to the efficient chain packing due to strong intermolecular hydrogen bonding and intermolecular n- n stacking interactions. Flowever, its unique macromolecular characteristics, rigid and straight heterocyclic molecular chain conformation as well as extensive hydrogen bonding also create challenges to fabricate it as asymmetric membranes with a desired morphology. When modified with coagulant and/or small molecules, it was found that finger-like macrovoids are formed inside the polymer layer, which weakens the membrane. Further, the swelling of membranes when in use need to be accounted for. It was found that when PBI is used, the ionic interactions between the electron-donating -N- groups in PBI and the electron-withdrawing group of the macrocycle having a particular size (in particular the electron-withdrawing sulfonic acid -SO3H groups in SCA8) may facilitate a uniform distribution of the macrocycles and firmly anchor them inside the PBI matrix without adversely affecting the desirable properties of the membrane. This allows industrial-size hollow fiber membrane modules for nanofiltration in organic solvents to be fabricated.

In some embodiments, a weight ratio of macrocycle to polymer is about 0.1 wt% to about 5 wt%. In other embodiments, the weight ratio is about 0.1 wt% to about 4 wt%, about 0.1 wt% to about 3 wt%, about 0.1 wt% to about 2 wt%, or about 0.1 wt% to about 1 wt%.

In some embodiments, the membrane composite has a Victoria blue B rejection of at least about 90% when the macrocycle loading is about 0.1 wt%. In other embodiments, the Victoria blue B rejection is at least about 92%, about 94%, about 96%, about 98%, or about 99%.

In some embodiments, the membrane composite has a methylene blue rejection of at least about 80% when the macrocycle loading is about 0.1 wt%. In other embodiments, the methylene blue rejection is at least about 80%, about 85%, about 90%, about 92%, about 94%, about 96%, about 98%, or about 99%.

In some embodiments, the membrane composite has a paracetamol rejection of at least about 30% when the macrocycle loading is about 0.1 wt%. In other embodiments, the paracetamol rejection is at least about 35%, about 40%, about 45%, about 50%, about 55%, or about 60%.

In some embodiments, the membrane composite has a pure methanol permeance of about 1 LMH/bar to about 1.8 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure methanol permeance is about 1.1 LMH/bar to about 1.8 LMH/bar, about 1.2 LMH/bar to about 1.8 LMH/bar, about 1.3 LMH/bar to about 1.8 LMH/bar, about 1.4 LMH/bar to about 1.8 LMH/bar, about 1.5 LMH/bar to about 1.8 LMH/bar, or about 1.6 LMH/bar to about 1.8 LMH/bar.

In some embodiments, the membrane composite has a pure acetonitrile permeance of about 1 LMH/bar to about 1.28 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure acetonitrile permeance is about 1.1 LMH/bar to about 1.28 LMH/bar, or about 1.2 LMH/bar to about 1.28 LMH/bar.

In some embodiments, the membrane composite has a pure acetone permeance of about 0.4 LMH/bar to about 1 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure acetone permeance is about 0.5 LMH/bar to about 1 LMH/bar, about 0.6 LMH/bar to about 1 LMH/bar, about 0.7 LMH/bar to about 1 LMH/bar, about 0.8 LMH/bar to about 1 LMH/bar, or about 0.9 LMH/bar to about 1 LMH/bar.

In some embodiments, the membrane composite has a pure ethanol permeance of about 0.48 LMH/bar to about 0.76 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure ethanol permeance is about 0.5 LMH/bar to about 0.76 LMH/bar, about 0.55 LMH/bar to about 0.76 LMH/bar, about 0.6 LMH/bar to about 0.76 LMH/bar, or about 0.7 LMH/bar to about 0.76 LMH/bar.

In some embodiments, the membrane composite has a pure ethyl acetate permeance of about 0.24 LMH/bar to about 0.36 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure ethyl acetate permeance is about 0.26 LMH/bar to about 0.36 LMH/bar, about 0.28 LMH/bar to about 0.36 LMH/bar, about 0.3 LMH/bar to about 0.36 LMH/bar, about 0.32 LMH/bar to about 0.36 LMH/bar, or about 0.34 LMH/bar to about 0.36 LMH/bar.

In some embodiments, the membrane composite has a pure tetrahydrofuran permeance of about 0.12 LMH/bar to about 0.24 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure tetrahydrofuran permeance is about 0.14 LMH/bar to about 0.24 LMH/bar, about 0.16 LMH/bar to about 0.24 LMH/bar, about 0.18 LMH/bar to about 0.24 LMH/bar, about 0.2 LMH/bar to about 0.24 LMH/bar, or about 0.22 LMH/bar to about 0.24 LMH/bar.

In some embodiments, the membrane composite has a pure toluene permeance of about 0.01 LMH/bar to about 0.08 LMH/bar when the macrocycle loading is about 0.3 wt%. In other embodiments, the pure toluene permeance of about 0.02 LMH/bar to about 0.08 LMH/bar, about 0.03 LMH/bar to about 0.08 LMH/bar, about 0.04 LMH/bar to about 0.08 LMH/bar, about 0.05 LMH/bar to about 0.08 LMH/bar, about 0.06 LMH/bar to about 0.08 LMH/bar, or about 0.07 LMH/bar to about 0.08 LMH/bar.

In some embodiments, the membrane composite has a stability in a methylene blue/methanol mixture for at least 7 days.

In some embodiments, the membrane composite has a A/,/V-dimethyl-4-nitroaniline (DMNA)/methylene blue (MB) separation factor of about 14.5.

In some embodiments, the membrane composite comprises: a) an outer layer of less than about 150 nm; and b) an inner layer comprising a bi-continuous sponge-like substructure. In some embodiments, the outer layer has a nodular substructure.

In some embodiments, when the membrane composite is a hollow fibre membrane, the membrane composite further comprises a porous inner surface.

In some embodiments, the membrane composite is characterised by a tensile strain at maximum elongation of about 40% to about 50%.

In some embodiments, the membrane composite is characterised by a maximum tensile stress of about 20 MPa to about 40 MPa, about 20 MPa to about 38 MPa, about 20 MPa to about 36 MPa, about 20 MPa to about 34 MPa, about 20 MPa to about 32 MPa, or about 20 MPa to about 30 MPa.

In some embodiments, the membrane composite is characterised by a Young's modulus of about 500 MPa to about 600 MPa, about 500 MPa to about 580 MPa, about 500 MPa to about 560 MPa, about 500 MPa to about 550 MPa, or about 500 MPa to about 540 MPa.

In some embodiments, the membrane composite is characterised by a pure water permeance of about 1 L m² h ¹ bar¹ to about 1.2 L rrr² h ¹ bar¹, about 1 L m² h ¹ bar ¹ to about 1.15 L m² h^-1 bar¹, or about 1 L m² h^-1 bar¹ to about 1.1 L m² h^-1 bar¹.

In some embodiments, the membrane composite is characterised by a MWCO of about 200 Da to about 300 Da, about 200 Da to about 280 Da, about 200 Da to about 260 Da, or about 200 Da to about 240 Da.

It was found that chemically crosslinking PBI by DBX tightens the pore size and its distribution, while the ionic interaction induced by SCA8 impregnation further sharpens the pore size and its distribution. As a result, PBI-DBX-SCA8 HFM has a small effective mean pore size and the sharp pore size distribution suitable for precise separation at an angstrom level.

In some embodiments, when the membrane composite comprises PBI and SCA8, its pure solvent permeance is characterised by an inverse of a solvent's viscosity. In some embodiments, the membrane composite is characterised by a Remazol Brilliant Blue rejection of at least about 90%, about 92%, about 94%, about 96%, about 98%, or about 99%.

In some embodiments, the membrane composite is characterised by a oil rejection of a 5% oil/acetone solution of more than about 90%, about 92%, about 94%, about 96%, about 98%, or about 99%.

In some embodiments, n is an integer selected from 5 to 8.

In some embodiments, the method further comprises a step of crosslinking the layer of polymer with a crosslinker having at least 2 pendant groups, the pendant group selected from aminoalkyl, haloalkyl, and a combination thereof. In other embodiments, the crosslinker is an amine crosslinker selected from 1,6-hexanediamine (HDA), tris(2- aminoethyljamine (TAEA) or a combination thereof.

In some embodiments, the crosslinker solution has a crosslinker concentration of about 5 wt%. In other embodiments, the concentration is about 1 wt% to about 10 wt%, about 1 wt% to about 9 wt%, about 1 wt% to about 8 wt%, about 1 wt% to about 7 wt%, or about 1 wt% to about 6 wt%. In some embodiments, the crosslinking step is performed for at least 12 h. In other embodiments, the crosslinking step is performed for at least 24 h. In some embodiments, the macrocycle is homogenous dispersed across the thickness of the layer of polymer.

In some embodiments, the macrocycle is characterised by a pore size of about 4 A to about 9 A, about 4.5 A to about 9 A, about 5 A to about 9 A, about 5.5 A to about 9 A, about 6 A to about 9 A, about 6.5 A to about 9 A, about 7 A to about 9 A, or about 7.5 A to about 9 A.

In some embodiments, the macrocycle comprises monomer residues of 5 to 12, 6 to 12, 7 to 12, 8 to 12, or 9 to 12. In some embodiments, the first polar moiety is selected from cyano, acyl, oxyacyl, acyloxy, amino, acylamino, aminoacyl, amidoximyl, oximyl, hydrazonyl, iminyl, hydroxyl, pyrrolyl, oxazolyl, imidazolyl, pyrazolyl, triazolyl, and a combination thereof. In some embodiments, the first polar moiety is selected from acyl, amino, acylamino, aminoacyl, and a combination thereof.

In some embodiments, the second polar moiety is selected from hydroxyl, sulfo, sulfonyl, amino, acyl, oxyacyl or a combination thereof.

In some embodiments, the membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; and wherein the macrocycle is sulfocalix[8]arene.

In some embodiments, the membrane composite, comprising: a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the polymer is polyamide-imide crosslinked with amine crosslinker and/or polybenzimidazole crosslinked with xylene; and wherein the macrocycle is sulfocalix[8]arene.

Details of the embodiments and workings of the invention are laid out below. In the embodiments that follows, the invention is described in relation to some conditions for consistency to showcase the present invention in different applications. However, the skilled person would understand that the invention is not limited to such.

Molecularlv homogeneous 5CA4 infiltration into the microporous polymer scaffold In some embodiments, amidoxime functionalized polymer of intrinsic microporosity PIM-1 (AOPIM) was used to provide a microporous scaffold for infiltrating SCA4 molecules because of several key advantages. Firstly, AOPIM polymers possess both N- containing amide and O-containing hydroxyl groups which are prone to forming hydrogen-bond or ionic interactions. Also, these functional groups can easily find highly similar counterparts, interaction-wise, in many other polymer backbones containing or easily functionalizable with N- or 0- groups. Secondly, the highly permeable but insufficiently selective nature of PIM-1 membranes could allow better visualization of the size-sieving effect of SCA4 and reduce the resulting permeability measuring errors for slow larger gases. Also, multiple important gas pairs have been shown to be separable using PIM-based membranes so that we can identify the breadth of applications of our PFI design.

By immersing the as-cast AOPIM1 films in methanol solutions of SCA4 molecules with varied concentrations, both the infiltration of SCA4 molecules into the microporous structure and the solvent exchange between methanol and the occluded casting solvent could occur simultaneously (Figure 1), which yielded the resultant SCA4-infiltrated AOPIM 1 (AOPIM1-SCA4) membranes. The degree of infiltration, defined as the actual mole ratio of SCA4 molecules to the polymer content in the membrane, was quantified to be 0.99%, 2.44%, 3.48% and 4.96%, respectively (simplified in the abbreviated sample name as 1%, 2%, 3% and 5%). Figure 1 shows a schematic for this facile infiltration process in addition to photos of the as-fabricated AOPIM1-SCA4 membranes. In comparison to prior techniques, conventional physically dispersed frameworks and polymers before membrane fabrication tend to form nanoscopic defects due to crystallization propensity, whereas protic-solvent-soluble macrocyclic molecules enable a molecular infiltration route into the already fabricated membrane. Figure 2 shows confirmation of the molecularly homogeneous infiltration of SCA4 molecules, in which (a) is photos of all AOPIM1-SCA4 samples showing no visible change in the physical appearances and apparent bendability; scanning electron microscopy (SEM) images and the sulfur element (red markers) distribution across (b) the surface and (c) the cross- section of the representative AOPIM l-SCA4-3% sample. From the physical appearances shown in the photos of the samples (Figure 2a), there were no visible changes of the color, texture, transparency or bendability found in the samples after SCA4 infiltration, indicating minimal disruption to the polymer chain packing. As shown in the imaging results from a scanning electron microscope coupled with an energy dispersive X-ray spectrometer (SEM-EDX), both the surface and cross-section of the representative AOPIM l-SCA4-3% sample (Figure 2b and c) appeared homogeneous and smooth and the elemental scans revealed homogeneous distribution of SCA4 molecules across the entire membrane containing the microporous scaffold. The line scan for the sulfur (S) element reflecting a relatively even distribution of the SCA4 concentration across the central region of the membrane also proved the ability of SCA4 to infiltrate deeply into the center of a 26 mm-thick dense film. In addition, the SCA4 molecules tend to repel one another (i.e. low self-assembly/aggregation propensity) after the proton transfer interactions with AO groups on AOPIM1 because they bear similar negative charges, and their distribution tends to follow the inherently fine distribution of AO groups in the membrane because of the strong interactions between them, which would promote the achievement of SCA4 infiltration with molecular homogeneity.

SCA4-polvmer interactions via hydrogen and ionic bonding

Figure 3 provides identification of polymer-SCA4 interactions and the bridging effect on polymer chains; (a) ¹H NMR spectra enlarged at the proton signals of the -OH on oxime groups; X-ray photoelectron spectra (XPS) for the N Is element of (b) AOPIM1 and (c) AOPIM-SCA4-3% membranes; (d) X ray diffraction (XRD) spectra to illustrate the evolution of interchain spacing (dotted arrows to show the trend, 0-5% indicated the degree of SCA4 infiltration); (e) distributions of PALS lifetimes and the corresponding free volume radii.

Bearing an amide ( = CR-NH2) and a hydroxyl-containing oxime (-CR]N-OH) functional group, the AO moieties could actually engage with the upper rim sulfonic pendant groups of SCA4 via different interactions, namely hydrogen-bonding and proton transfer interactions (i.e. ionic bonding). In the enlarged ¹H NMR spectra (Figure 3a), the triplet signal assigned to the hydroxyl proton on oxime groups experienced an appreciable decrease in both peak intensity and sharpness as the degree of infiltration increased. This was most probably caused by these protons' involvement in extensive hydrogen bonding with the numerous oxygen atoms on sulfonic groups. In contrast, for a control sample incorporated with 5 mol% calix[4]arene-25,26,27,28-tetrol (referred to as CA4t and its membrane sample was abbreviated as AOPIMl-M-CA4t-5%), the aforementioned triplet signal remained largely unaffected. Since CA4t is a SCA4 analog without sulfonic or any pendant groups, this spectral difference verified the necessity of sulfonic groups in forming extensive hydrogen bonding.

On the other hand, the amide group with a highly proton accepting nitrogen (N) tended to undergo proton-transfer interactions with the sulfonic moieties with high protonating tendency. This was evidenced by a much stronger N⁺ signal (402.25 eV) in the N Is X- ray photoelectron spectra (XPS) of an AOPIMl-SCA4-3% membrane sample as compared with a pristine AOPIM1 sample (Figure 3b and c). In addition, the signal peak of =CR-NH2 (amide) shifted and increased from a binding energy of 400.50 eV to 400.90 eV, which could be attributed to the electron withdrawal from its N atom during the formation of ionic bonds. It is believed that SCA4 can be potentially applied to other polymers.

The bridging effect of SCA4 on polymer packing

With four sulfonic groups on the upper rim pointing outwards in four directions, each SCA4 molecule could simultaneously interact with multiple nearby chains or segments on the polymer backbones via extensive hydrogen and ionic bonding, resulting in a bridging effect within the membrane's microporous scaffold. As revealed by X-ray diffraction (XRD) patterns of AOPIM1 membranes (Figure 3d), two broad amorphous peaks with respective d-spacings of 6.39 A and 4.90 Å were observed, signifying the characteristic existence of two major types of microporous structures in PIM-1 membranes with different average inter-chain (or inter-segmental) distance. The former accounted for the less efficiently packed regions arising from the contorted ladder-type polymer backbones that formed larger micropores, while the latter represented the more densely packed regions from which the ultrafine micropores were derived. Due to the interaction-induced bridging effect, both d-spacing values appeared to experience a slight decreasing trend after the membrane was infiltrated with SCA4 molecules, but the largely rigid 3D structure of SCA4 could help limit the tightening or contractive effect on the overall polymer chain and inter-chain spacing, which was most probably the reason why these d-spacing changes were relatively very small.

It was also observed that the larger microporous regions tended to undergo slightly more contraction (i.e. from 6.39 to 6.20 A) than the ultrafine ones (i.e. from 4.90 to 4.86 A) as the degree of infiltration increased, and the fact that the former consisted of more spacious and readily available micropores for accommodating the infiltrating SCA4 molecules than the latter should provide a probable explanation for that. Interestingly, both contractions seemed to slow down significantly after 2% infiltration as reflected by the observation that the average d spacing gradually stopped evolving after this, which could occur after all spacious regions have been internally bridged by an identical 'space holder'.

Enhanced mechanical stability from SCA4's conformational flexibility The bridging effect between SCA4 molecules and nearby polymer chains helped strengthen the chain-to-chain adhesion forces, hence improving the overall elastic modulus of AOPIM1-SCA4 membranes. As shown in Table 1, 1% to 3% infiltration of SCA4 endowed the membranes with greater tensile strength than that of the pristine one. These reinforced mechanical properties may also arise from the fact that SCA4 could distort its 3D conformation, to a certain extent, from a more bowl-like to a more cup-like conformation (illustrated by the distorted SCA4 bowl-shapes in Figure 5) to adapt to the compressive forces exerted by strained micropores during sample deformation. As a closed-loop organic polymer molecule with methylene bridges between the phenol sub-units, SCA4 molecules actually possessed some conformational flexibility on top of their overall structural rigidity, which also resulted in their bottom opening exhibiting a small range of sizes around a mean value of 3.0 A. The key benefit of such conformational flexibility is that it could help minimize any localized mechanical weak spots that could compromise the membrane's pliability. Although the overall improving trend of the mechanical properties halted at 5% infiltration most likely due to the excessive occupancy of micropores that deprived SCA4 molecules of enough space to adaptively maneuver their shapes, the mechanical robustness of AOPIM1- SCA4-5% was still as good as that of the pristine membrane, revealing the potential of approaches for tackling vigorous practical applications.

Table 1. Membrane bulk density, mechanical strength and insoluble content.

Sharper pore size distribution in AOPIM1-SCA4 membranes

The average pore sizes, their distribution and the total fractional free volume (FFV) of AOPIM1-SCA4 membranes were measured and analyzed by positron annihilation lifetime spectroscopy (PALS) and the numerical results are summarized in Table 2 and the plotted distribution is shown in Figure 3e. It was observed that the intensity and the mean free volume radii (i.e. pore size) of both the larger (14, r4) and ultrafine (13, r3) micropores appeared to experience a mild decrease as the degree of SCA4 infiltration increased, which could be a reasonable trend because hydrogen-bond and ionic bridging effects should be able to mildly tighten the overall polymer packing. The total FFV also appeared to demonstrate a slightly decreasing trend, which could be the result of both these bridging effects as well as the occupancy of micropores by SCA4 molecules. Although the magnitude of these numerical changes seemed rather small, they might still be able to reflect the sufficiently altered nanoporosity within a membrane that could possibly give rise to huge gas transport property changes as corroborated by other studies with small changes in the PALS results. This was also one of the reasons for employing the PALS technique which could ensure high precision. Table 2. PALS lifetimes and corresponding free volume radii obtained at ambient temperature and pressure.

It is noteworthy that the pore size distribution illustrated a distinctive difference between the evolutions of these two types of free volume elements. As the degree of infiltration increased, the extent of size heterogeneity was clearly reduced in the larger microporous regions with their pore size distribution becoming sharper and narrower, while the ultrafine pores decreased in their mean size without manifesting obviously narrower distribution. This disparity in the evolution of the pore size distribution revealed the preferential lodging of SCA4 into the larger microporous regions because such an enhancement in size homogeneity was most possibly caused by their favorable accommodation of these identical 'space holders' that size-standardized the originally broadly sized micropores. In contrast, the lodging of SCA4 into the ultrafine microporous regions would be more sterically unpreferable so that only a reduction in mean size was observed due to the overall bridging effect being transduced from larger microporous regions via the networked microporous structure of PIM-1 without allowing its distribution to be homogenized.

Permeation tests for multiple industrially relevant gases

Figure 4 shows gas transport parameters and sorption behaviors of AOPIM1-SCA4 membranes; (a) Permeability and selectivity changes as a function of the degree of SCA4 infiltration; (b) correlation between the permeabilities of various gases and their kinetic diameters; (c) sorption isotherms with data points fitted to the dual-mode sorption model for N2, CPU and CO2 pure gases obtained at 35 °C and a pressure up to 10 bar; (d) gas solubility (cm³ (STP) per cm³ membrane bar) and diffusivity (xlO⁷ cm² s ¹) as a function of the degree of SCA4 infiltration for N2, CPU and CO2 at 35 °C and 3.5 bar.

The pure-gas permeation tests on AOPIM1-SCA4 membranes for PI2, O2, N2, CPU and CO2 were carried out at 35 °C and 3.5 bar and the results are summarized in Table 3. By analyzing the gas permeability and selectivity as a function of the degree of SC4 infiltration (Figure 4a), a general trend of decreasing permeability for all the gases was observed, which was reasonably expected because of the previously observed decreasing trend of the total FFV. Flowever, gases with different sizes displayed vastly different trends. H2 permeability consistently remained within the same order of magnitude with the least extent of reduction experienced throughout the whole studied range, while CO2 permeability experienced a moderately higher decrease due to its relatively larger size than H2 as well as its cylindrical shape which could result in more difficult passage at lateral positions. In contrast, larger-sized gas molecules, like N2 and CFU, experienced a one order of magnitude or higher permeability drop, especially at higher degrees of infiltration, and this asynchronism of permeability changes brought about very interesting evolution of gas selectivity.

Table 3. Summary of permeability and selectivity for H2, O2, N2, CFI4 and CO2 at 35 °C and 3.5 bar.

At 1% infiltration, there was about 65% and 73% increase in the H2/N2 and FU/CFU selectivity, respectively, with almost unaffected H2 permeability because the partial coverage over the entire microporous scaffold by SCA4 cavities started to exhibit some but not complete size-discrimination against the larger N2 and CFU gas molecules. Once the degree of infiltration reached 2%, the size-sieving effect became drastically more prominent with the FI2/N2 and H2/CH4 selectivity becoming 4.6 and 10.6 times, respectively, as high as that of the pristine membrane, but at the expense of merely 15.7% FI2 permeability. A 3.3-fold enhancement in CO2/CFU selectivity was also observed and despite the moderate loss of CO2 permeability, the tremendous improvement in selectivity over CH4 still boosted the overall performance to well transcend the recent upper bound (shown in Figure 6d). More interestingly, after the SCA4 infiltration, the AOPIM1 membranes even transformed from being originally CO2/H2-selective to increasingly H2/CO2-selective due to the rise of molecular-sieving diffusion within the membranes and a 3.5 times higher H2/CO2 selectivity was obtained even at 2% infiltration. Being corroborated by the evolution of pore size, SCA4 molecules tended to settle down in the spacious but less- or non-size-selective larger micropores during solution infiltration. Subsequently, they could act as molecular gatekeepers that guarded these non-selective regions within the membrane from indiscriminate gas passage and also redirected gas transport within the entire interconnected microporous network as elucidated by an illustrative diagram shown in Figure 5. Figure 5 illustrates the molecular gatekeeping mechanism in the AOPIM1 microporous scaffold. The size-sieving 3D cavity of SCA4 preferentially grants passage to small H2 and CO2 while strongly impeding large N2 and CPU. Since the full coverage of all non-selective regions would require a certain degree of infiltration to be reached, this explained the dramatic performance jump from 1% to 2% infiltration because there should exist a degree of saturated infiltration between 1% and 2% at which all non- selective micropores are remodelled with size-sieving properties. Beyond saturation, further infiltration of SCA4 molecules could result in more unnecessary transport resistance against gas diffusion than additional molecular-sieving effectiveness so that selectivity improvement slowed down from 2% onwards. This phenomenon actually has an analogous explanation in the defect healing process in thin film membranes if we consider non-selective micropores to be nano-defects in the selective layer, and the infiltration of and ionic bridging effect by SCA4 actually 'healed' them while bringing in a new gatekeeping cavity. Once all the 'defects' were healed, selectivity enhancement slowed down even with additional infiltration. This so-called 'defect-healing' phenomenon could probably also explain the observation that the evolution of size distribution of the larger micropores became less significant after around 2% infiltration (Figure 3e) as additional SCA4 would become ineffective in further homogenizing these micropores after they are fully infiltrated with the same SCA4 'space holders'.

The emergence of effective molecular-sieving structures in these membranes consequently resulted in a strong negative correlation between the pure-gas permeability and the gas kinetic diameter (Figure 4b), especially at and beyond the degree of saturated infiltration. In comparison, those of PIM-1 and AOPIM1 membranes exhibited a much weaker size dependence because of the presence of non-selective microporous regions without being guarded by molecular gatekeepers. It is important to note that the SCA4 cavity possesses a small range of size around 3.0 A owing to its conformational flexibility so that it might not perfectly refuse the passage of gas molecules with a kinetic diameter bigger than 3.0 A. Nevertheless, a strong molecular- sieving nature was still achieved as both the likelihood and the rate of diffusion of larger gases were tremendously lowered as their size increased. For example, C02 (3.3 A) was granted passage much more preferentially than CH4 (3.8 A) for enabling high CO2/CH4 selectivity.

Dual-mode sorption analyses and aas transport properties

From the sorption isotherms of AOPIM1-SCA4 (Figure 4c) samples obtained with N2, CFU and CO2 respectively at 35 °C over a range of pressures from 0 to 10 bar, there were no significant changes found in the overall gas solubility after the infiltration of SCA4 for these three gases, which reflected that the chemistry and microporous structure of AOPIM1 polymers remained largely unaltered. These isotherms displayed a good fit to the dual mode sorption model of glassy polymers, and those of the solid SCA4 powder also resembled dual-mode behaviors, demonstrating SCA4's inherent affinity for gases. By examining the fitted dual-mode capacity parameters summarized in Table 4, the change of sorption sites for each probe gas was explored so that the proposed guarded microporous network structure within AOPIM1-SCA4 membranes could be further elucidated. For N2 and CFU that have larger sizes than CO2, the number of Langmuir sites decreased while the corresponding Flenry's law sites increased as signified by the corresponding lower C'H but higher kD constant values, respectively, after the infiltration of SCA4. These opposing changes in the two sorption sites revealed that the original Langmuir sites, contributed by the excess free volume elements in the AOPIM1 membranes, became less accessible to N2 and CFU after they were guarded by the size-discriminative SCA4 gatekeepers. As a result, these sites started to resemble the Flenry's law sites for N2 and CFU. In addition, the capture ability of SCA4 cavities also contributed to the increased affinity parameter, b, for N2 and CFU. In contrast, being the smallest and most condensable gas among the three tested gases, CO2 experienced minimal changes in the nature of its corresponding sorption sites as the kD and C'H constants did not show any significant variation after SCA4 infiltration and only some capture effect was observed as the parameter b slightly increased, which also supported the preferential passage of CO2 through SCA4 cavities over larger gases. Interestingly, kD, C'H and b values were found to remain roughly unchanged from a certain degree of infiltration onwards which was again between 1% and 2%, similar to the degree of saturated infiltration. A probable explanation is that additional SCA4 molecules could only generate a marginal impact on the overall affinity of these sites towards N2 and ChU. However, the mechanistic understanding of this phenomenon might require a separate study and was not further explored here.

Table 4. Dual-mode sorption fitting constants for N2, CH4 and CO2 at 35 °C and a pressure up to 10 bar.

From the sorption isotherms, the solubility (S) of each gas at 3.5 bar was obtained and is summarized in Table 5 together with their diffusivity (D) determined from the solution-diffusion correlation (permeability (P) = D x S). As illustrated by the evolution of these two gas transport parameters (Figure 4d), the solubility of all three gases (N2, CH4 and CO2) remained mostly constant as the degree of SCA4 infiltration increased, whereas their diffusivity dropped more significantly with CH4 experiencing the most drastic reduction. This showed that the solubility selectivity of H2 against each of these three gases as well as that of CO2 against CH4 remained almost unchanged while their diffusivity selectivity was greatly enhanced. Therefore, the dramatic enhancements in H2/N2, H2/CH4 and CO2/CH4 selectivity were contributed primarily by the diffusivity selectivity increase arising from the size-sieving SCA4 cavities that favored the diffusion of smaller gases. The previously discussed C02/H2-to-H2/CC>2 selectivity transition was also a result of this domination of diffusivity selectivity over solubility selectivity after the introduction of SCA4 molecular gatekeepers as the originally non-selective microporous regions gradually evolved to a highly size-sieving state. Table 5. Summary of gas transport parameters, including solubility (measured from sorption isotherms) and diffusivity for N2, CFU and CO2 at 35 °C and 3.5 bar

Control experiments illustrating SCA4's gatekeeping role and unblocked cavities

Two control experiments were carried out to reveal the gas passage through and the unblocking of SCA4 cavities. For the former, AOPIM1 membranes were infiltrated with inorganic sulfuric acid via a similar procedure to crosslink their microporous structure, which yielded AOPIMI-H2SO4 membranes. Under the same infiltration conditions, AOPIMI-H2SO4 membranes showed a higher degree of crosslinking as reflected by the larger amount of insoluble content in their samples (Table 1) as well as the slightly higher permeability but much lower selectivity than those of the analogous AOPIM1- SCA4 membranes. If the passage of smaller gas molecules did not occur through SCA4 cavities and whatever permeability and selectivity changes were solely due to the bridging effect on the polymer packing, by virtue of SCA4's larger size than that of sulfuric acid, the AOPIM1-SCA4 membranes logically should have displayed higher permeability but lower selectivity than AOPIMI-H2SO4, especially when H2SO4 molecules were actually able to produce a higher degree of crosslinking among the polymers. Yet, the opposite was observed and exceptionally high selectivity for multiple gas separations was obtained only using AOPIM1-SCA4 membranes, indicating that the gas passage occurred through a highly size-sieving window which in our control cases could only be the SCA4 open cavity. The molecular gatekeeping role of the SCA4 open cavity in the polymer microporous scaffold was also evidenced from this. As for the latter, SCA4 molecules were physically blended with AOPIM1 polymers to fabricate the AOPIM1-M- SCA4 membranes as a method control sample and all of their separation performances fell far below those of the AOPIM1-SCA4 membranes with no strong molecular-sieving nature being manifested. This was most possibly caused by the blockage or filling of the SCA4's single cavity by polymer chains during the intense physical mixing process. In contrast, solution in filtrating SCA4 in the already solidified membranes where polymer chains exhibited much less mobility than in a solution state could minimize the likelihood of cavity blockage, which is one key advantage of the PFI design that is not realizable by conventional inorganic or framework-type nanomaterials.

Overall separation performances and their stability against aging Figure 6 shows pure-gas separation performances of AOPIM1-SCA4 membranes compared with recent upper bounds and other high-performance polymer-based membranes. Plots of (a) H2/N2, (b) H2/CFU, (c) FI2/CO2 and (d) CO2/CFU separation performances at 35 °C and 3.5 bar with either 2008 or 2015 or both upper bounds (0- 5% denotes the degree of infiltration, the numbers in parentheses indicate days of aging, the numbers at the end of abbreviated names of literature studies indicate the pyrolysis temperature in °C, and arrows are drawn as a guide to the eye).

The drastic enhancements in H2/N2 and H2/CF-I4 selectivity with small compromises in H2 permeability drove the AOPIM1-SCA4 membranes from 2% onwards to perform much closer to or even well above the state-of-the-art 2015 upper bounds recently updated for these two separations (Figure 6a and b) and to also be comparable with or even outperform many other top performance polymer-based membranes, including newly synthesized PIMs, thermally treated or crosslinked membranes and some mixed-matrix membranes (MMMs). Similarly, the CO2/CFI4 and FI2/CO2 performances were also boosted to well transcend 2008 upper bounds and became comparable to other state- of-the-art reports (Figure 6c and d).

In addition, the AOPIM1-SCA4 membranes demonstrated good performance stability against physical aging as all these pure-gas separation performances at 2% and 3% infiltration either moved even closer to or consistently stayed well above the 2008 or 2015 upper bounds after 60 days (Figure 6a-d). During aging, some losses of gas permeability were inevitable because the residual insufficiently selective regions closed down and the ultrafine micropores also shrank in size due to the gradually relaxed polymer chains collapsing excess free volume. Nevertheless, both these changes not only resulted in a more size discriminative overall pore structure, but might also amplify the prominence of the gatekeeping effect of SCA4 cavities, leading to more than proportionate gains in the H2 selectivity. CO2/CFU selectivity did not undergo further enhancement after aging due to the collapsed excess free volume that retarded its solubility selectivity, but its overall aged performance remained persistently beyond the upper bound while that of the pristine membrane fell below. Meanwhile, the A0PIM1- SCA4 membranes also showed attractive O2/N2 separation performance which was high above the 2008 upper bound and well maintained after aging, suggesting the potential of the present invention for air purification applications as well. The mixed-gas H2/CH4 performance of A0PIM1-SCA4 membranes tested using an equimolar gas mixture (Table 6) was lower than the pure-gas performance with respect to both permeability and selectivity, which is typical because of the competitive sorption among gas species in a mixed permeation environment. Nonetheless, their mixed-gas H2/CH4 selectivity still sharply increased with the degree of infiltration (up to 24 times as high) while the maximal reduction in H2 permeability was only 35%. As a result, the overall mixed gas separation performance was able to transcend the 2008 pure-gas upper bound easily.

Importantly, to obtain such ultra-efficient, durable and industrially relevant multiple gas separations, the membrane composite is highly facile and potentially scalable, and can strategically overcome the longstanding issues in conventional composite membranes, especially those prepared by the often problematic physical-mixing route. Also, huge separation performance enhancement could be achieved with only a small consumption of SCA4 (2-3 mol%), suggesting an extremely high enhancement-to-consumption leverage that was uncommon in conventional composite membranes which typically required 10-30 wt% of incorporated porous agents to exhibit some advantageous effects, if any. Lastly, although the illustrated size-sieving around 3.0 Å already successfully targeted a variety of energy and environmentally important gas separations, it is believe that modulation by size and functionality tuning via base size change or lower rim substitution coupled with different types of water-soluble upper rims can possibly bring this technology over to many other polymers that target an even wider range of molecular separations.

Conclusions

The present invention discloses an unconventional post-fabrication infiltration (PFI) membrane to incorporate external intrinsically porous agents into polymer membranes after dense film formation. In this membrane composite, issues of interfacial void and pore blockage, which are often exacerbated by the conventional physical-mixing strategy, can be overcome by choosing a water-soluble organic macrocyclic molecule, in particular, 4-sulfocalix[4]arene (SCA4). Its unique feature of coexisting solubility in common membrane-treating protic solvents, like methanol, and the presence of a 3D intrinsic open cavity in its molecule-sized body enabled SCA4 to homogeneously infiltrate the microporous scaffold contained in the amidoxime-functionalized PIM-1 (AOPIM1) membranes via the formation of extensive hydrogen or ionic bonding, which could offer generalizable design principles across other existing polymers. The AOPIM1- SCA4 membranes fabricated based on this ultra-facile PFI design demonstrated drastically enhanced molecular-sieving characteristics with outstanding separation performances for multiple important gas pairs being achieved, including FI2/CO2, H2/N2, H2/CH4, CO2/CFI4 and O2/N2, which well surpass the 2008 or 2015 upper bounds and are as good as or even better than those of many other polymer-based membranes with outstanding performances. Investigation into the pore size distribution, sorption behavior, gas transport properties and relevant control experiments mechanistically revealed the role of SCA4 molecules as size-sieving molecular gatekeepers guarding non-selective regions within the microporous network in AOPIM1 membranes. Being backed by a large pool of size-tunable and functionalizable water-soluble macrocyclic molecules, the present membrane composite offers potential applicability in an excitingly wide array of energy-intensive molecular applications.

As disclosed herein, calix[n]arene is a type of close-loop macrocyclic molecules, containing several repeating units of phenolic blocks with a hollow cup-like or frustum structure. As illustrated in Figure 7a, b, the wide upper-rim, narrow lower-rim, and the number of repeating units (n) can be molecularly tuned to synthesize various kinds of calix [n]arene. Without wanting to be bound by theory, the inventors further believe that the molecular- tunability and molecular-sieving characteristics of calix[n]arene can be utilised for making membrane composites by ion ica lly bonding calix[n]arene such as sulfothiacalix[4]arene (STCAss) and sulfocalix[4]arene (SCA) in a polymer network (such as polyamide), and examine their potential for OSFO. Thiacalix[n]arene is a subgroup of calix[n]arene where the methylene bridge between each phenolic block is replaced by sulfur atoms, which may alter the lower-rim dimension and ultimately the size-sieving properties. Since both STCAss and SCA have multifunctionalities, they may interact with the polyamide layer closely and affect its formation and separation performance. Herein, it is shown that (1) functionalization of the polyamide network can allow incorporation of close-loop size-selective macrocyclic molecules, and (2) simultaneously improvement of the permeability and selectivity of the membrane composites for OSFO processes can be obtained.

Surface morphologies

Figure 8a presents the surface morphology of the pristine membrane. The pristine polyamide layer has a typical ridge-and-valley morphology with relatively big "leaves". After incorporating STCAss and SCA (Figure 8b and 8c), the resultant polyamide layers still possess ridge-and-valley morphologies but the "leaves" become smaller. In addition, the thickness of the polyamide layer decreases with the addition of macrocycles possibly due to the acidic nature of STCAss and SCA, which possess four sulfonate and sulfonic acid groups in one molecule, respectively. Among them, TFN-SCA-1.5 has the thinnest selective layer. The pH values of the m-phenylenediamine (MPD) solutions confirm our hypothesis. The MPD solutions containing 1.5 wt% STCAss and SCA have pH values of ~6.6 and 5.5, respectively, while the pristine MPD solution has a pH value of ~9. Since the sulfonate and sulfonic acid groups in STCAss and SCA may interact with MPD, their presence interferes the interfacial polymerization process. Moreover, the proton dissociated from the sulfonic acid groups would inhibit the HCI generation during the interfacial polymerization because it shifts the reaction equilibrium backward. Both factors result in a lower degree of cross-linking reaction between MPD and trimesoyl chloride (TMC) during the interfacial polymerization, leading to a less crumpled polyamide layer.

A comparison of the surface morphology between TFN membranes embedded with STCAss and SCA indicates that TFN-STCAss-1.5 has smaller"leaves" than TFN-SCA-1.5, although the latter has a slightly lower pH than the former. This may be caused by the Na⁺ ions released from STCAss. The interfacial polymerization reaction generally takes place in two steps: (i) MPD in the aqueous phase diffuses toward the interface and reacts with TMC to form a nascent polyamide layer without the ridge-and-valley structure; (ii) Marangoni convection further prompts the migration of MPD vigorously to react with TMC, thus pushing and bending the nascent polyamide layer and forming a ridge-and-valley structure. Since Marangoni convection is affected by the surface tension near the water/hexane interface and the addition of inorganic electrolytes, such as Na⁺ and K⁺, elevates the surface tension of water, the presence of Na⁺ from STCAss in the MPD solution may increase the surface tension near the water/hexane interface, and reduce the Marangoni convection, leading to a smaller "leaves".

The surface topology of these polyamide layers was further probed by atomic force microscopy (AFM) to determine the surface roughness. As shown in Figure 8, the order of surface roughness is consistent with the observation from field emission scanning electron microscopy (FESEM) images, which follows TFC-0 > TFN-SCA-1.5 > TFN- STCAss-1.5. This order further confirms the effects of pH and Na⁺ on the formation of polyamide layer during the interfacial polymerization.

Surface chemistries

The surface chemistries of the pristine and modified polyamide layers were characterized by Fourier transform infrared spectroscopy (FTIR) under the attenuated total reflectance (ATR) mode and X-ray photoelectron spectroscopy (XPS) using freestanding polyamide films. Figure 9 displays the ATR-FTIR spectra of TFC-0, TFN-STCAss- 1.5, and TFN-SCA-1.5. The pristine polyamide layer consists of the typical N-H and C-0 bonds of amide groups, which show sharp stretching peaks at 3300 cm^-1 and 1620 cm^-1, respectively. In comparison, after STCAss and SCA incorporation, the sharp N-H stretching peak of the amide group is covered by a broad O-H stretching peak in the range of 3200-3650 cm^-1, while an S=0 stretching peak at 1149 cm^-1 appears. The O- H and S=0 bonds belong to the hydroxyl and sulfonate/sulfonic acid groups of STCAss and SCA, respectively. This indicates the successful incorporation of both STCAss and SCA into the polyamide layers. However, it is difficult to distinguish the difference between TFN-STCAss-1.5 and TFN-SCA-1.5 from the ATR-FTIR spectra because the C- S bond shows a weak peak in FTIR that is usually covered by other peaks. Therefore, XPS is employed to further investigate the surface composition and bonding information of the free-standing membrane composites (TFC and TFN).

Table 7 tabulates the surface compositions of these polyamide layers. The S atom can hardly be detected in the pristine polyamide layer although there is a tiny amount of sodium dodecyl sulfate (SDS; 0.2 wt%) in the MPD aqueous solution. By incorporating STCAss and SCA, the S contents in the polyamide layers increase to 0.34 At% and 0.26 At%, respectively. The S content of TFN-STCAss-1.5 is higher than that of TFNSCA-1.5 because each STCAss molecule consists of 8 S atoms and each SCA molecule possesses only 4 S atoms. The C contents are similar for the three polyamide layers; however, the O contents increase while the N contents decrease after the incorporation of STCAss and SCA. This is due to the fact that STCAss and SCA possess high concentrations of O elements, owing to both hydroxyl and sulfonate/sulfonic acid groups but no N elements.

To further analyze the surface chemistry, the N 1 s and S 2 p XPS spectra are deconvoluted and displayed in Figure 10a. For the N peak, there are three fitted peaks; namely, (1) primary amine from the partially cross-linked MPD monomer, (2) secondary amine for the polyamide group, and (3) quaternary amine. The percentages of their peak areas are calculated and listed in Table 8.

Table 8. Peak area percentages of N Is for TFC-0, TFN-STCAssl.5 and TFNSCA-1.5.

The percentage of the secondary amine decreases significantly with the incorporation of STCAss and SCA, while the contents of both primary and quaternary amines increase slightly. The decrease in secondary amine and the increase in primary amine further confirm the lower cross-linking degrees in these two modified polyamide layers as observed in FESEM. Moreover, the increase in quaternary amine content may be due to the formation of ionic bonding between the sulfonate/sulfonic acid groups and unreacted amine/amide groups. As elucidated in Figure 7a, b, the sulfonate and sulfonic acid groups in STCAss and SCA are strong electron withdrawing groups; meanwhile, the amine and amide groups tend to donate electrons to them so that ionic bonds are formed, as depicted in Figure 10b. This ionic bonding not only ensures strong interactions between the incorporated macrocycles and the polyamide network, but also facilitates the dispersion of STCAss and SCA nano-fillers inside the polyamide layer. As a result, the scanning electron microscopy-energy dispersive X-ray spectroscopy (SEM- EDX) image of S elements in TFN-STCAss-1.5 shows an even distribution across the membrane surface, indicating a homogenous dispersion of STCAss in the polyamide network. The ionic bonding between the polyamide network and STCAss/SCA would make the developed TFN membranes superior to those conventional TFN membranes in terms of stability and homogeneity because the latter tends to only have physical interactions between the polyamide layer and nano-fillers.

For the S 2 p XPS results, there is no obvious peak for the pristine polyamide layer, which is consistent with the surface composition analysis. In contrast, the polyamide layers embedded with STCAss and SCA show two separate peaks and one sharp peak, respectively. The two peaks for TFN-STCAss-1.5 are attributed to C-SO3 (168.5 eV) and C-S-C (164 eV), since there are two types of sulfur atoms in STCAss. For TFN-SCA-1.5, since SCA only contains C-SO3, there is only one sharp peak observed at 168.5 eV. Thus, by using XPS, TFN-STCAss-1.5, and TFN-SCA-1.5 can be simply distinguished from their S 2 p spectra.

Membrane microstructures

In order to investigate the depth profile of microstructures and the evolution of free volume and thickness in polyamide layers, TFC-0, TFN-STCAss-1.5, and TFN-SCA-1.5 were examined by positron annihilation spectroscopy (PAS). As illustrated in Figure 11, the S- and R- parameters are used to characterize the changes of free volume and micro-voids as a function of positron penetration depth (i.e., membrane thickness), respectively. The S-parameter represents the intensity of free volume, where a larger S-parameter stands for increased free volume cavities and/or a higher free volume. Similarly, the R-parameter describes the pore size distribution, where a larger R- parameter means that the voids (in nm to pm sizes) become larger and/or their quantities increase.

As shown in Figure 11a, b (enlarged one), the S-parameters of all membranes increase sharply first, which is caused by the back diffusion and scattering of positroniums near the membrane surface. Then, TFC-0 exhibits a different trend from the other two. Its S-parameter fluctuates at a certain value and then increases further. The initial fluctuation represents the dense polyamide layer and the later increase indicates the transition from the dense polyamide layer to the cross-linked polyimide substrate. As the free volume intensity inside the membrane is not ideally homogenous along the membrane depth, the S-parameter normally fluctuates around a certain level instead of a smooth line. In contrast, the S-parameters of TFN-STCAss-1.5 and TFNSCA-1.5 reach the maximum values first at ~1-1.2 KeV and then decrease at higher incident energy, matching the S-parameter of the substrate in TFC-0 at ~2.5 KeV. Thus, the S-parameter corresponding to the positron energy of 0.5-2.5 KeV represents the polyamide layer. To facilitate easy comparison, eye-guiding lines that represent the average values of S- parameters in the range of 0.5-2.5 KeV, have been added as dotted lines in Fig. 5b. The average values of S-parameters, corresponding to the polyamide layers, follow the order of TFC-0 < TFN-STCAss-1.5 ~ TFN-SCA-1.5. In addition, there is hardly any overlap between TFC-0 and TFN -STCAss- 1.5/TFN -SCA- 1.5 in the range of 0.5-2.5 KeV, so the changes in free volume intensity by incorporating STCAss and SCA are considered to be significant. This confirms that the incorporation of STCAss and SCA into the polyamide network can significantly increase its free volume intensity.

The R-parameters of all membranes present a similar trend, as plotted in Figure 11c, d. The R-parameter decreases at the beginning and reaches the lowest value, followed by a drastic increase, which forms a "V-shape" curve. The initial decrease and subsequent increase imply the existence of a thin dense-selective layer. Thus, the distance between the initial point and the bottom of the "V-shape" can be interpreted as the dense-layer thickness, and the minimum value of the "V-shape" indicates the lowest intensity of voids. Consistent with the observation from FESEM images, the dense-layer thickness follows a descending order of TFC-0 > TFN-STCAss-1.5 > TFN-SCA-1.5. In addition, the pristine TFC membrane has the smallest intensity of voids, while the other two TFN membranes possess a similar intensity of voids. In summary, the pristine polyamide layer has the lowest intensity of free volume and voids as well as the largest thickness; the STCAss incorporated polyamide layer is slightly thinner and has a relatively larger intensity of free volume and voids; while the SCA-modified polyamide layer possesses similar intensity of free volume and voids as the STCAss-incorporated polyamide layer, but has the thinnest polyamide layer among them.

OSFO performance and transport properties

The OSFO performances of TFC-0, TFN-STCAss-1.5, and TFN-SCA-1.5 were quantified under FO and pressure-retarded osmosis (PRO) modes, using pure ethanol and 2M LiCI in ethanol as feed and draw solutions, respectively. Figure 12a, b shows their ethanol flux (Jw), reverse solute flux (J_s), as well as J_s/Jw. The pristine TFC membrane has a relatively low ethanol flux and a high reverse salt flux under FO and PRO modes, which are not desirable for an FO membrane. In contrast, the STCAss-incorporated membrane has a dramatic increase in ethanol flux under both FO and PRO modes due to its thinner polyamide layer and a larger free volume, as validated by PAS. Surprisingly, TFN-SCA- 1.5 has almost the same ethanol flux as TFC-0, although the former has a higher free volume and a thinner polyamide layer than the latter, as revealed by PAS. Meanwhile, the reverse solute flux follows the order of TFC-0 > TFN-STCAss-1.5 > TFN-SCA-1.5. In other words, the STCAss and SCA-incorporated membranes have much lower reverse solute fluxes than the pristine one, which also does not follow the trend of their free volumes observed by PAS. These interesting phenomena may arise from the fact that the close-loop macrocyclic molecules, i.e., STCAss and SCA, not only function as pore formers but also act as precise molecular sieves that only allow small molecules (i.e., smaller than their cavity sizes) to pass through and reject or block the big ones. According to the molecular structures of STCAss and SCA drawn and optimized by Material Studio^®, Figure 7 reveals that the diameters of the lower cavities in STCAss and SCA are ~5.1 A and 4.2 A, respectively. Since the kinetic diameter of an ethanol molecule is ~4.4 A, ethanol can easily pass through the small opening of STCAss but may not go through SCA. As a consequence, even though TFN-SCA-1.5 has a higher free volume and a thinner polyamide, it exhibits a similar ethanol flux to the pristine TFC membrane. Similarly, the small cavities of both STCAss and SCA may reject Li⁺ and Cl- ions in ethanol although their solvated diameters in ethanol are not known. Nevertheless, the hydrated diameters of Li⁺ and Cl- are 7.64 A and 6.64 A, respectively, which are much larger than the small cavity sizes of STCAss and SCA. Flence, the STCAss- and SCA incorporated membranes reject Li⁺, Cl-, and LiCI, and show low reverse solute fluxes in OSFO tests.

In order to confirm our hypothesis, the transport properties of these membranes were investigated by measuring their pure ethanol permeances (A) and salt rejections (Rs) under the deadend filtration mode. Figure 12c shows that TFN-STCAss-1.5 possesses the highest pure ethanol permeance, while TFN-SCA-1.5 has the lowest pure ethanol permeance. The rejection toward LiCI, NaCI, and KCI follows the order of TFN-SCA-1.5 > TFN-STCAss-1.5 > TFC-0, which reconfirms our hypothesis that both small cavities of STCAss and SCA can effectively reject Li⁺, Cl-, and LiCI. For all three membranes, the rejections follow the order of LiCI > NaCI > KCI, due to the different sizes of solvated ions in ethanol. In addition, the transport properties of ethanol measured under the dead-end filtration mode are consistent with those under OSFO modes. TFN-STCAss- 1.5 always has the highest ethanol flux among these membranes. Although TFC-0 has a slightly higher pure ethanol permeance than TFN-SCA-1.5 under the dead-end filtration mode, their ethanol fluxes under OSFO modes are similar. This may be due to the fact that TFC-0 has a higher reverse solute flux. Thus, the driving force across this membrane is reduced that results in a lower ethanol flux under OSFO tests.

The solvent transport mechanism for the developed membranes was also investigated by measuring pure solvent permeances of water, methanol, ethanol, and then calculating their d iff usivities in the membranes according to the solution-diffusion model. As shown in Figure 12d, the solvent diffusivities of all membranes decrease as the solvent molecules become larger. This is in accordance to the diffusivity correlating equations, i.e., diffusivity is inversely related to molecular size. However, the trend of TFNSCA-1.5 is slightly different from those of TFC-0 and TFN-STCAss-1.5. The former has a much lower ethanol diffusivity than the latter because it has a stronger size- sieving effect to block ethanol, as explained previously. The tunable molecular-sieving properties of the fabricated membrane composites were also observed in aqueous solutions, when ethylene glycol (EG), diethylene glycol (DEG), and glucose were selected as probing solutes for aqueous filtration tests (Figure 13). Interestingly, the rejection of TFN-SCA-1.5 toward EG is much larger than those of TFC-0 and TFN- STCAss-1.5. However, the rejections of both TFN-SCA-1.5 and TFN-STCAss-1.5 toward DEG are similar and are only slightly higher than that of TFC-0. Since the diameter of EG is 4.7 A, it may only be excluded by SCA but not by STCAss. However, the diameter of DEG is 5.82 A, which could be rejected by both SCA and STCAss. Meanwhile, a similar trend in rejections can also be observed for glucose, which is attributed to its diameter of >7 A. This phenomenon further implies that both STCAss and SCA can molecularly tune the molecular-sieving properties of the fabricated TFNs not only in ethanol but also in aqueous solutions.

Optimization of STCAss loading

Based on the above results, STCAss is selected as the suitable nano-filler for further studies because it can increase free volume of the polyamide layer, reduce its thickness, and maintain high solute rejection. Thus, the STCAss loading in the MPD solution was varied from 0 to 2.0 wt%, in order to investigate the optimal loading and obtain the best OSFO performance. Tables 9 and 10, and Figure 14 summarize AFM, XPS, and FESEM results, respectively.

Table 9. Surface roughnesses of the pristine TFC and TFNs with different STCAss loadings.

Generally, an increase in STCAss loading in the MPD solution results in the polyamide surface with a higher S content and a smoother polyamide layer due to the aforementioned acidic effect. The smooth polyamide layer may be desirable to minimize fouling and maintain a sustainable ethanol flux.

Figure 15 illustrates the OSFO performance under FO and PRO modes as a function of STCAss loading. In both modes, the ethanol flux rapidly increases while the reverse solute flux significantly decreases, with an increase in STCAss loading from 0 wt% to 1.5 wt%, leading to a desirable Js/Jw value. The high ethanol flux guarantees a high ethanol recovery while the low reverse solute flux helps in maintaining the driving force across the membrane and ensuring a stable ethanol flux. In addition, the low Js/Jw value indicates a low leakage of the draw solute to the feed stream and minimizes the potential hazard caused by the reverse flux of draw solutes. However, the ethanol flux drops and the reverse solute flux bounces back significantly, when an excess amount of STCAss (2 wt%) is embedded in the polyamide layer. This is due to defect formation as evidenced by the increase in surface roughness and the appearance of bulges (Table 9, Figure 14). The TFN formed by incorporating 1.5 wt% STCAss in the MPD solution results in the best OSFO performance. Comparing to the pristine TFC membrane, its ethanol fluxes are almost doubled under both FO and PRO mode. Interestingly, the percentage of increment under the FO mode (93.7%) is slightly lower than that of the PRO mode (112.2%) because of the ICP effect. The former has a lower Js/Jw value of 0.07 than the latter of 0.12 using 2M LiCI as the draw solution, respectively.

Discussion

The polyamide network has been functionalised with two calix[n]arene, STCAss and SCA for OSFO by incorporating them into the MPD monomer solution via interfacial polymerization. Under FO and PRO modes, the optimized TFN membrane containing 1.5 wt% STCAss in the MPD solution exhibits higher ethanol fluxes (FO: 3.38 Lrrr² h^_1 (LMH) and PRO: 5.37 LMH) and lower reverse solute fluxes (FO: 0.23 grrr² h^_1 (gMH) and PRO: 0.64 gMH) than the pristine TFC membrane (FO: 1.74 LMH, 1.42 gMH and PRO: 2.53 LMH, 5.41 gMH). FESEM, AFM, and PAS results confirm the synergetic effects of STCAss on the polyamide network due to: (i) the thinner and smoother polyamide layer caused by the acidic nature of STCAss, (ii) the ionic bonding formed between the sulfonate groups of STCAss and the amide/amine groups of the polyamide network, (iii) the increased free volume in the polyamide network, and (iv) the size-exclusion effect from the hollow cup-like structure of STCAss. As the TFN-STCAss-1.5 exhibits the optimal OSFO performance, its applicability to simultaneously concentrate pharmaceuticals and reclaim organic solvents is further tested using two feed solutions of 2 g L^_1 paracetamol (MW = 151.2 g mole^-1) and tetracycline (MW = 444.4 g mole^-1) in ethanol and 2M LiCI as the draw solutions. The membrane exhibits stable ethanol fluxes of 3.41 LMH and 3.26 LMH, with rejections of 96.1% and 99.6% toward paracetamol and tetracycline, respectively. To compare with OSN or low-pressure organic solvent reverse osmosis (OSRO), filtration tests were conducted using 0.2 g L^_1 paracetamol and tetracycline solutions as feed solutions for TFN-STCAss-1.5. As shown in Table 11, the results clearly suggest that the incorporation of STCAss into the polyamide layer has great potential to molecularly design TFN membranes for OSFO, OSN, and low-pressure OSRO, in order to simultaneously concentrate pharmaceutical products and recover organic solvents. The present invention shows that TFN membranes when fabricated with macrocyclic molecules, such as STCAss, can simultaneously improve permeability and selectivity for OSFO. In addition, the calix[n]arene-functionalized polyamide network may also have great potential for sea water desalination, waste water treatment (e.g., heavy metal and organic pollutants removal), solvent reclamation, and drug purification via pressure-driven process.

The inventors have also applied the present invention to hollow fiber membranes. This is predicated on the motivatation to overcome the problems of hollow fiber membranes by modifying integrally skinned asymmetric membranes with much enhanced molecular sieving capability.

To this end, the inventors endevoured to use macrocycles in the fabrication of hollow fiber membranes. Macrocycles refer to a group of cyclic molecules containing several membered rings, such as cyclodextrins, calixarenes and cucurbiturils. Similar to metal organic frameworks and covalent organic frameworks, macrocycles possess intrinsic cavities that could size-selectively reject solutes which are spatially larger than their cavities. Moreover, some of them have functional groups that can be reacted or utilized for further modifications. They can also have intrinsic sieving properties which can be employed to enhance the size-selectivity of the membranes. For instance, cyclodextrins have been utilized to react with acyl chlorides and form polyester membranes with high permeances and size-selectivity. As an example, a polyamide-imide polymer, Torlon®, was chosen as the polymer material to fabricate the integrally skinned asymmetric HFM substrate, due to its good chemical and mechanical properties. To further improve its chemical stability, the as- spun HFM was crosslinked with a tripodal amino-crosslinker, tris(2-amino)ethylamine (TAEA). To enhance its molecular sieving capability, 4-sulfocalix[8]arene (SCA8) which possesses 8 phenolic blocks in a close loop membered ring, was impregnated into the crosslinked HFM. The ionic interactions between the electron-withdrawing sulfonic acid groups of SCA8 and the electron-donating amine groups of TAEA may facilitate SCA8 to be uniformly distributed and firmly anchored inside the membrane. Thus, it is shown that integrally skinned asymmetric HFMs can be revitalize with enhanced size-selective molecular separation for OSN applications by impregnating them with SCA8. It is believed that this is the first attempt to invigorate crosslinked asymmetric HFMs for OSN with superior sieving capability.

Membrane morphology

The morphology of the as-spun HFM was observed by a field emission scanning electron microscope (FESEM), as displayed in Figure 16. Its outer surface is relatively dense with a few cracks resulted from the drying and cryo-fracturing processes during sample preparation. In contrast, its inner surface possesses (1) a lot of pores because the bore fluid contains mostly NMP and (2) polymer particles due to the phase inversion from a lean polymer phase. Macrovoids are distributed just beneath the outer selective layer owing to the water (external coagulant) intrusion to the nascent fiber during the phase inversion. The enlarged outer section contains a dense selective layer of ~917 nm, while the inner section possesses a loose structure full of big cells to facilitate solvent transportation. In addition, the membrane morphology does not apparently change after HDA and TAEA modifications, as well as SCA8 impregnation, as shown in Figure 17.

Membrane chemistry

Figure 18 provides the chemical bonding information of the as-spun, HDA-XIinked and TAEA-XIinked HFMs as detected by attenuated total reflectance- Fourier transform infrared (ATR- FTIR) spectroscopy. The pristine polyimide-amide polymer, Torlon^®, contains both imide and amide groups, which have characteristic peaks at 1777 cm ¹ and 1712 cm ¹ for imide C=0, 1635 cm ¹ for amide C=0, 1537 cm ¹ for amide C-N and 1369 cm ¹ for imide C-N. Meanwhile, both HDA and TAEA successfully crosslink Torlon^® by converting its imide groups into amide groups as evidenced by the weakened imide C=0 and C-N peaks in the ATR-FTIR spectra of both HDA- and TAEA-XIinked HFMs. Flowever, the chemical stability of FIDA-XIinked HFMs are not sufficiently improved because they become gelatinous after immersion in N,N-dimethylformamide (DMF) for 6 months. In contrast, TAEA-XIinked HFMs still remain intact. This implies that TAEA can result in crosslinked Torlon^® HFMs with stronger chemical stability than HDA, possibly because the former has tripodal amines while the latter has dipodal amines. Figure 18 also reveals that SCA8 has been successfully anchored onto the TAEA-crosslinked HFM via

The elemental compositions of the HFMs were analyzed using XPS, as tabulated in Table 12. The N content in the pristine HFM is relatively low (6.32%), which gradually increases to 9.02% and 11.07% after crosslinking with HDA and TAEA, respectively. Compared to the HDA-XIinked HFM, the higher N content of the TAEA-XIinked HFM can be ascribed to the higher N content per TAEA molecule than that per HDA one. As there is no N element in SCA8, the increasing S content, O/N and S/N ratios can be interpreted as an increasing SCA8 amount impregnated inside the hollow fibers. Figure 19 further elucidates the bonding information of the prepared HFMs by deconvolution of the N Is peak. The as-spun HFM is made of a polyimide-amide polymer, which contains only secondary (-NH-) and tertiary (-N<) amines. Thus, only one peak is needed to fit the spectrum. After being crosslinked with HDA and TAEA, the primary amine (-NH2) and quaternary amine (NH3⁺) peaks appear. The primary amine originates from the partially crosslinked HDA or TAEA, while the quaternary amine arises from the protonated amine groups. Table 13 tabulates their percentages based on the peak areas. Interestingly, the TAEA-XIinked HFM has a slightly larger percentage of the primary amine than the HDA-XIinked HFM, because the three amines in TAEA have a lower possibility to be fully reacted than the two amines in HDA. This also arises from the fact that the amine groups in HDA are much stronger nucleophiles than those of TAEA because the former has higher pKa values than the latter (i.e., 9.83 and 10.93 vs. 8.56, 9.59 and 10.29). After impregnating SCA8 into the HFM, the percentage of the quaternary amine peak increases from 20.56% to 27.62%. Moreover, the percentages of both primary and secondary/tertiary amines decline slightly because SCA8 may have been ionically bonded to either the primary or tertiary amines of TAEA on the HFM. Since the sulfonate groups of SCA8 tend to withdraw electrons from the electron-sufficient groups of both primary and tertiary amines of TAEA, an ionic interaction is established between TAEA and SCA8 such that TAEA provides the anchoring site for SCA8 to firmly attach to the membrane matrix. Figure 20 schematically illustrates the TAEA induced crosslinking mechanism and the ionic interaction with SCA8. Elemental mapping via FESEM-EDX shows that both the S element (an indicator of SCA8) and the N element (an indicator of amines) have similar patterns, and they are homogenously distributed across the cross-section. These evidence that the ionic interaction between TAEA and SCA8 may facilitate the homogenous impregnation of SCA8 molecules inside the HFM.

Table 13. Percentages of peaks in N Is deconvolution

Percentage (%)

Primary

Secondary/Tertiary amine Quaternary amine amine

As-spun HFM 100

HDA-XIinked HFM 24.80 56.12 19.08

TAEA-XIinked

28.21 51.23 20.56

HFM

TAEA-SCA8-0.3 23.19 49.19 27.62 Membrane microstructure

In order to study the effects of crosslinking and SCA8 impregnation on microstructure of HFMs, the as-spun HFM, FI DA- and TAEA-XIinked HFMs, as well as TAEA-SCA8-0.3 were characterized by PAS. S-parameter, which represents the change of free volume intensity along the membrane depth, was measured. Generally, a larger S-parameter designates a higher intensity of free volume, where the free volume cavity becomes more and/or larger. Flowever, the S-parameter is also influenced by the quenching effect of a polyimide polymer that traps positrons and restricts positronium formation, resulting in a low S-parameter. As shown in Figure 21a, the as-spun FIFM has the lowest S-parameter among all fabricated FIFMs, possibly due to a low free volume intensity and the quenching effect. By crosslinking with FIDA and TAEA, the imide rings in Torlon^® are converted to amide groups, which significantly alleviates the quenching effect and results in larger S-parameters for FIDA- and TAEA-XIinked FIFMs. In addition, the enlarged S-parameters may be also caused by the increase in free volume, as evidenced by pure water permeance (PWP) and pore size distribution, which will be discussed in the later section. The TAEA-XIinked FIFM has a slightly higher S-parameter profile than the FIDA-XIinked FIFM possibly because TAEA is slightly bigger than FIDA in terms of molar volume. Thus, the former creates a slightly bigger free volume cavity than the latter. After SCA8 impregnation, the S-parameter of TAEA-SCA8-0.3 decreases significantly as compared to that of the original TAEA-XIinked FIFM. This decrease implies that SCA8 diffuses and attaches to the polymer matrix, not only occupying the free volumes among polymer chains but also tightening the structure via ionic interaction.

Figure 21b displays another parameter, R-parameter, which evaluates the evolution of voids in the range of nm-pm, along the membrane depth. Similar to S-parameter, a larger R-parameter denotes larger and/or more voids inside the membranes. All R- parameters for the prepared membranes decrease first to the lowest level and then increase, forming a "V-shape" curve. This trend is caused by the presence of a thin and dense selective layer, so the width of this valley can be interpreted as the thickness of the selective layer. All the prepared membranes have similar valleys in terms of width and depth because (1) they were made or crosslinked from the same integrally skinned FIFMs and (2) they have almost the same thicknesses of the selective layers as displayed in Figure 16 and 17. For the region just beneath the selective layer, the R-parameter shows a similar trend to S-parameter, where TAEA-XIinked FIFM has the highest R- parameter, followed by FIDA-XIinked FIFM, TAEA-SCA8-0.3, and the as-spun one. This reveals that the crosslinking modifications induced by HDA and TAEA will enlarge the pores inside the HFMs, while the impregnation of HFMs with SCA8 will take up some space of the voids.

OSN separation performance Chemical crosslinkina with FIDA and TAEA

To investigate the effects of chemical crosslinking with FIDA and TAEA on pore size distribution, the fabricated FIFMs were tested using aqueous feeds with different PEG molecular weights. As displayed in Figure 22a, the as-spun FIFM has a PWP of 8.7 L nr ² h ¹ bar¹ and an MWCO of 933 Da. In contrast, the FIDA-XIinked FIFM has a PWP of 26.3 L m² h^-1 bar¹, which is almost triple of the as-spun one. Similarly, its MWCO jumps to 4559 Da due to the increases in free volume and pore size as demonstrated by PAS. A similar phenomenon has also been observed previously during the crosslinking modification of a polyimide flat-sheet membrane (36). Interestingly, the PWP value of the TAEA-XIinked FIFM is further increased to 45.6 L m² h^-1 bar¹, despite its MWCO is only slightly larger than the FIDA-XIinked FIFM.

Figure 22b discloses the pore size distributions of the as-spun, FIDA- and TAEA- crosslinked FIFMs. The as-spun FIFM has the sharpest distribution and smallest mean effective pore size, but the pore size distribution curve shifts to the right and becomes broad after it is crosslinked with either FIDA orTAEA. The FIDA-XIinked FIFM has a slightly smaller mean effective pore size than the TAEA-XIinked one, but the former has a broader pore size distribution than the latter. This interesting phenomenon may arise from different molecular sizes and structures of FIDA and TAEA. FIDA is a slender and linear molecule (8.8 A x 2.5 A) with dipodal amines. It is spatially more compact than TAEA, but its long flexible alkane chains make the chain packing and folding easier and more flexible. In comparison, TAEA has a spatially bulky structure (7.8 A x 4.3 A) with tripodal amines. Its tripodal amines provide more active sites to effectively crosslink the membrane and sterically restrict the chain movement. Since the molecular size and structure counterbalance with each other, both FIDA- and TAEA-crosslinked FIFMs have similar MWCOs.

Figure 22c compares the performance of the as-spun, FIDA- and TAEA-XIinked FIFMs to separate rose bengal (RB, MW= 1017.6 g mol ¹) from a methanol solution. Compared with the aqueous system (Figure 22a), the methanol permeance is much lower than the water permeance. Consistent with the trends of PWPs and effective mean pore size, the as-spun HFM has the smallest pure methanol permeance of about 1 L m² h^-1 bar¹ with an RB rejection of 99.9%. The HDA-XIinked HFM has a higher pure methanol permeance of 4.64 L m² h ¹ bar¹ but its RB rejection decreases to 87.05%. The TAEA-XIinked HFM has the highest methanol permeance of 9.24 L m² h ¹ bar¹; however, its RB rejection drops to 66.5%. Therefore, a further modification of the HFMs is necessary to enhance their molecular sieving properties.

Impregnation of HFMs with SCA8

The effect of SCA8 loading on separation performance of HFMs was investigated in methanol. Three dye molecules; namely, paracetamol (size: 7.8 Ax 4.2 A, MW= 151.2 g mol ¹), methylene blue (MB, size: 12.1 Ax 4.7 A, MW= 319.9 g mol ¹) and Victoria blue B (VBB, size: 12.7 A x 11.3 A, MW= 506.1 g mol ¹) were employed as probing solutes. As shown in Figure 23a, the pure methanol permeance declines dramatically by increasing the SCA8 loading because it can diffuse into the membrane matrix and alter its free volume and voids, as revealed by PAS. Consequently, the paracetamol rejection increases gradually as a function of SCA8 loading. However, the MB rejection increases extraordinarily from 23.2% to 85.2% upon impregnation with only 0.1 wt% SCA8. This dramatic increase in rejection can be ascribed to the enhanced sieving capability enabled by SCA8. As SCA8 has a cavity size of about 9.2 A, paracetamol with a length of 7.8 A, could still pass through it. However, MB, with a length of 12.1 A, tends to be blocked by SCA8. Thus, the TAEA-SCA8-0.1 has a rejection difference between MB and paracetamol more than two-folds of TAEA-XIinked HFM. Moreover, the rejections of a much larger dye, VBB, are all greater than 95% for SCA8 impregnated HFMs because VBB has both length and width larger than lnm. Clearly, the SCA8 impregnated HFM is promising to precisely reject solutes with a size larger than 1 nm.

To understand the fundamentals of solvent transport across the SCA8 impregnated HFM, the permeances of 7 organic solvents, containing polar protic, polar aprotic and nonpolar solvents, were measured using TAEA-SCA8-0.3 (Figure 23b). Methanol has the highest permeance, followed by acetonitrile, acetone, ethanol, ethyl acetate, tetrahydrofuran and toluene. Generally, a solvent with a small molecular size, low viscosity and higher affinity towards the membrane usually exhibits a high permeance across the membrane. Figure 23c fits the measured permeances as a function of solubility parameter (S), viscosity (m) and molar volume (V) of the solvent. A linear correlation can be established between the permeance and a combination of δm^-1V^-1.

Figure 23d presents the evolution of permeance and rejection as a function of time by conducting a long-term performance test of TAEA-SCA8-0.3 in an MB/methanol mixture for 7 days and collecting the permeance and rejection every day. The membrane exhibits a steady permeance and a consistent rejection during the entire 7-day period, indicating that the SCA8 impregnated membrane has very stable structural integrity. These results confirm the good stability of the TAEA-SCA8 modified FIFM in methanol and its potential for industrial applications.

Mixed solutes separation and benchmark

In order to demonstrate the precise molecular sieving capability of TAEA-SCA8-0.3 FIFM, it was utilized to separate a mixture consisting of two small solutes; namely, N,N- dimethyl-4-nitroaniline (DMNA, size: 6.9 A x 2.5 A, MW= 166.2 g mol ¹, color: yellow) and MB (size: 12.1 Ax 4.7 A, MW= 319.9 g mol^-1, color: blue). The mixture has a dark teal color in methanol. By filtrating it through the FIFM, the permeant shows a light green color, as illustrated in Figure 24a. The UV absorption spectra confirm that MB is almost completely rejected by the FIFM, while DMNA can mostly pass through the membrane. The separation factor, which measures the efficiency of the size sieving process, is about 14.5, signaling that the SCA8 impregnated FIFM has an excellent capability to separate DMNA and MB. Figure 24b illustrates the molecular sieving mechanism across the membrane. Since SCA8 has a cavity size of around 9.2 A, it can size-selectively reject molecules with a size larger than 1 nm (MB) but let the small ones (DMNA) pass. Table 14 compares the OSN performance between the TAEA-SCA8 modified FIFMs and other reported OSN FIFMs in the literature. The newly developed FIFMs have higher rejections with comparable permeances as compared to other OSN HFMs.

Discussion

In this work, an integrally skinned Torlon^® HFM has been fabricated and modified with TAEA and SCA8 to successfully enhance its chemical stability and molecular sieving capability for OSN. SCA8 is uniformly impregnated into the HFMs by a solvent infiltration method, attributable to the ionic interactions between SCA8 and TAEA. The SCA8 impregnated HFMs exhibit enhanced molecular sieving capability: the MB (size >lnm) rejection increases significantly after SCA8 impregnation; while, the paracetamol (size <lnm) rejection only increases gradually. Furthermore, the HFMs impregnated with 0.3wt% SCA8 can effectively separate a mixture of two solutes having close molecular weights, i.e., DMNA (MW= 166.2 g mol ¹) and MB (MW= 319.9 g mol ¹), with a separation factor of 14.5. It is believed that the present invention allows for integrally skinned FIFMs with sharp sieving capability of FIFMs to be applicable in sustainable separation processes.

Membrane morphology of PBI+5CA8 hollow fiber membrane

Figure 25a displays the morphologies of the as-spun PBI FIFMs. It consists of a selective dense layer less than 100 nm, a bi-continuous sponge-like substructure, and a porous inner surface. A nodular structure appears on the outer surface probably because the outer surface undergoes a metastable and nucleation growth state during the liquid- liquid demixing in the coagulation bath. Finger-like macrovoids are completely removed in the cross-section since relatively mild coagulants, IPA and 14/86 wt% DMAC/IPA, were applied outside and inside the nascent FIFMs respectively during the phase inversion. A sponge-like substructure and an inner surface with high porosity are formed owing to the delayed phase separation. The inter-connected porous structure can effectively reduce transport resistance and increase the solvent flux. Finger-like macrovoids within FIFMs are generally considered as undesirable weak points, they may lead to structural failure under high pressure and temperature operations. A macrovoid- free structure may effectively dissipate the external forces and guarantee adequate mechanical strength to resist high trans-membrane pressures. Moreover, the FIFM morphology does not show apparent changes after the DBX crosslinking modification and SCA8 impregnation, as observed in Figure 25b.

Membrane chemistry of PBI+SCA8 hollow fiber membrane

Qualitative bonding information of as-spun PBI, PBI-DBX and PBI-DBS-SCA8 FIFMs was measured by means of ATR-FTIR. The absorption peak from 2500 to 3600 cm^-1 of the as-spun PBI FIFMs can be attributed to the N-FI groups in imidazole rings, while the peaks at 1612, 1590, 1443, 1286 and 801 cm ¹ are associated with the imidazole and benzene rings and their conjugation in the as-spun PBI membrane. The broad band around 3160 cm^-1 indicates the stretching vibration of N-FI groups participated in hydrogen bonding. The peak centered at 3420 cm^-1 is related to the stretching vibration from the isolated non-hydrogen bonded N-FI group. For PBI-DBX FIFM, two distinctive peaks at 2920 cm ¹ and 2850 cm ¹ are presented. They are assigned to C-H (the terminal C of the DBX) stretch and C-N (chemical crosslinking between the PBI molecular backbone and DBX) stretch, proving the crosslinking reaction between the imidazole rings of PBI and DBX. The FTIR results also reveal that SCA8 has been effectively anchored into the PBI-DBX HFM matrix through ionically interactions because of the existence of S=0 bonds at the peak of 1040 cm ¹ from its ATR-FTIR spectra.

Mechanical properties of PBI+SCA8 hollow fiber membrane

The elongation at break, tensile stress, and Young's modulus were measured by an Instron tensile machine. As expected, the DBX crosslinked PBI hollow fiber possesses higher mechanical strength and tensile strain than the pristine one. The stiffness of the PBI FIFM is also enhanced due to the increased stress at break and Young's modulus from the covalent modification. Clearly, the DBX crosslinking modification has synergistically improved the overall mechanical properties of PBI membranes. A comparison of mechanical properties between PBI-DBX and PBI-DBX-SCA8 FIFMs indicates that the ionic interaction through SCA8 impregnation further enhances the overall mechanical properties but the improvements are smaller than the DBX crosslinking modification.

The tensile strain at maximum elongation of PBI-DBX-SCA8 FIFM is 43.12 ± 4.81 %, the maximum tensile stress is 27.85 ± 2.52 MPa, and the Young's modulus is 526.1 ± 23.1 MPa.

Membrane microstructure PBI+SCA8 hollow fiber membrane

To understand how DBX and SCA8 modify the microstructure of PBI membranes, pose sizes and pore size distributions of the fabricated PBI FIFMs were characterized using aqueous solutions containing PEGs with different molecular weights. The as-spun PBI FIFM has a PWP of 2.6 L m² h ¹ bar¹ and a MWCO of 1896 Da. After the DBX crosslinking, the PBI-DBX FIFM has a PWP of 1.23 L m² h ¹ bar¹ and its MWCO drops significantly to 340 Da owing to the decreased pore size. Interestingly, the MWCO of the PBI-DBX-SCA8 FIFM further decreases noticeably to 217 Da but its PWP is only slightly smaller than the PBI-DBX FIFM (1.05 L m² h ¹ bar¹). Figure 28 discloses their pore size distributions and explains the trends of their PWP and MWCO values. The as-spun PBI FIFM has the largest effective mean pore size and the broadest pore size distribution. The chemically crosslinking modification by DBX tightens the pore size and its distribution, while the ionic interaction induced by SCA8 impregnation further sharpens the pore size and its distribution. As a result, the PBI-DBX-SCA8 HFM has the smallest effective mean pore size and the sharpest pore size distribution suitable for precise separation at an angstrom level.

OSN separation performance: solvent transport properties

To investigate the solvent transport phenomena through the DBX crosslinked and SCA8 impregnated PBI HFMs, 10 organic solvents, consisting of polar protic, nonpolar and polar aprotic solvents, were applied to measure their permeances across PBI-DBX-SCA8 HFMs under 10 bar. Figure 29A shows the permeances of different solvents through PBI-DBX-SCA8 HFMs. Basically, pure solvent permeances follow the trends: (1) for polar protic solvents: acetone > MeCN> methanol> ethanol > IPA; (2) for nonpolar solvents: hexane > ethyl acetate > toluene; and (3) for polar aprotic solvents: THF > DMF. Moreover, the PBI-DBX-SCA8 HFMs remain intact and flexible in these 10 organic solvents and exhibit neglectable axial elongation except in DMF where an extension of 1.8% is observed. This neglectable axial elongation makes it feasible to fabricate PBI- DBX-SCA8 HFMs into industrial-size HFM modules.

Apart from the swelling, the relationships between membrane permeances and solvent properties (i.e., molar volume, molecular weight, solubility parameter and viscosity) have received great attention by the OSN community. In general, a solvent which has a lower viscosity, smaller molecular size and higher solvent-membrane affinity tends to display a higher permeance through the membrane. A well-fitted linear relationship of the obtained permeances as a function of molar volume (MV) and viscosity (n) of the solvents cannot be established between the permeance and the combined term of MV/h because R2 = 0.5325. Thus, another correlation is explored between solvent viscosity and permeance. Figure 29B shows a better correlation with R2 = 0.788, indicating an increasing permeance is linked with a decreasing viscosity. Therefore, the permeances of pure solvents follow the sequence: hexane > acetone > methanol > ethanol > IPA, whereas the solvent viscosities are in a reversed order. This implies that solvent viscosity is a major factor influencing solvent transport across PBI-DBX-SCA8 HFMs. Interestingly, hexane has the highest permeance, followed by ethyl acetate, acetone, acetonitrile, methanol, toluene although the as-spun PBI and DBX/SCA8 modified PBI HFMs are all hydrophilic. Therefore, the highest hexane permeance may result from its extremely low viscosity and very linear molecular shape. Separation performance

Three dyes with different molecular weights and molecular configurations are tested: N,N-dimethyl-4-nitroaniline, methylene orange and Remazol Brilliant Blue. The performance of the PBI-DBX-SCA8 HFM to separate DMNA, MO and RBB from their acetone solutions was compared. The membrane shows a 99.5 % rejection to RBB and a low rejection of 17.5 % to DMNA owing to the differences in their molecular weights. The membrane also exhibits a rejection of 77.5% to MO although it is expected to be over 90% based on its MWCO measured in water. This discrepancy may be due to the linear configuration of MO molecules with a size of 4.8 A xl4.9 Å and the membrane swelling induced by different solvents. Comparing with the pure solvent permeance of 2.2 L m² h ¹ bar¹, the acetone permeances drop slightly to 1.90, 1.85 and 1.95 L m² h ¹ bar¹ for solutions containing DMNA, MO and RBB, respectively. Separation of mixed solutes

To verify the precise molecular sieving potential of PBI-DBX-SCA8 HFMs, they were applied to segregate two solutes with different molecular weights from a mixture; namely, DMNA (MW= 166.2 g mol ¹, size: 6.9 A x 2.5 A, colour, light yellow) and RBB (MW= 626.5 g mol ¹, size: 11.9 Ax 15.8 A, colour: blue). The mixture solution bears a blue colour in acetone. The permeated effluent shows a light yellow colour when filtrating the mixture through the FIFMs. The UV absorption spectra confirms that the HFMs almost completely reject RBB, while mostly letting DMNA permeate through. The separation factor of DMNA over RBB through the membrane segregation is about 176. Therefore, the SCA8 impregnated PBI-DBX HFMs have an excellent capability to segregate DMNA and RBB. Because SCA8 has an open cavity size of around 9.2 A, it can enhance the spatial-selective rejection to a molecule with a size bigger than 11.9 A (RBB) while permitting smaller one (DMNA) permeate through.

Integrally skinned PBI hollow fiber membranes applied in the solvent recovery from oil/acetone mixture

The integrally skinned PBI HFMs were fabricated from a 18 wt% solution in dimethylacetamide (DMAc) through the dry-jet wet phase inversion process. It consisted of chemically crosslinking a polybenzimidazole (PBI, MW = 27,000 g mol ¹) hollow fiber membrane with a,a'-dibromo-p-xylene (DBX) and further modified the crosslinked PBI membranes with 4-sulfocalix[8]arene. The fully crosslinked PBI hollow fibers exhibit solvent stability in the polar and nonpolar organic solvent, especially in the aprotic polar solvents. The obtained PBI hollow fiber membranes were packaged to a one-inch module through epoxy potting.

Figure 26 exhibits the performance of HFM modules as a function of oil concentrations at 10 bar. The oil rejection decreases slightly from 99.08 % to 98.28%. The permeance decreases from 0.85 to 0.40 L m² h ¹ bar¹ when the oil concentration increases from 5 wt% to 20 wt%. These permeance variations are probably due to the osmotic pressure differences built up between the permeate and feed sides together with the severe concentration polarization at the feed side.

Figure 27 shows the evolution of solvent permeance and oil rejection as a function of time during a long-term performance test. The PBI-DBX-SCA8 FIFM module was operated in a 10wt% oil/acetone mixture under 10 bar for 30 days and the permeance was collected every day. The 1-inch FIFM module exhibits a steady permeance of around 0.35 L m² h ¹ bar¹ after 10 days and a continuously increased oil rejection > 99.0% during the entire 30-day period. Over a long term operation of 16 days under 10 bar, the PBI FIFM module showed an increased oil (triglyceride, MW 885g/mol) rejection up to 99.60%. This indicates that the DBX crosslinked and SCA8 impregnated PBI HFMs have very stable structural integrity and chemical stability. They may have great perspective for the large-scale industrial applications in organic solvents.

Examples

Materials

Monomers for the synthesis of polymers with intrinsic microporosity-1 (PIM-1), namely, 5,5',6,6'-tetrahydroxy-3,3,3,3'-tetramethyl-l,l'-spirobisindane (TTSBI, 97%) and 2,3,5,6-tetra-fluoroterephthalonitrile (TFTPN, 99%), were purchased from Alfa Aesar and Matrix Scientific, respectively. Prior to use, the former was purified via recrystallization from methanol and the latter was purified via vacuum sublimation. A hydroxylamine solution (50% aqueous) was purchased from Merck and used directly for the functionalization of PIM-1 without any pretreatment. 4-Sulfocalix[4]arene (SCA4, 97%, (C₇HS04S)4, 744.74 g mol ¹) and calix[4]arene-25,26,27,28-tetrol (CA4t, 95%, (C7H6O)4, 424.49 g mol ¹) were purchased from Merck and dehydrated at 120 °C under vacuum for one night before every use. N,N'-Dimethylformamide (DMF, 99.%) purchased from VWR Chemicals was purified via vacuum distillation at 60 °C before reaction and polymer dissolution. Anhydrous potassium carbonate (K2CO3, >99.5%) and sodium hydroxide pellets (NaOH, >98.5%) from Merck, methanol (MeOH, 99.8%) and ethanol (EtOH, HPLC, 99.8%) from Fisher Chemical, n-hexane (95%) from Tedia, tetrahydrofuran (THF, 99.7%) and hydrochloric acid (HCI, 32% aqueous) from VWR Chemicals, and concentrated sulfuric acid (FI2SO4, 95% aqueous) from Avantor® were directly used as received. All purified gases (>99.5%), including H2, O2, N2, CH4, and CO2, and also the mixed gas of 50% H2 and 50% CH4 were supplied by Air Liquide Pte. Ltd. (Singapore).

Polymer synthesis

The synthesis protocol for the PIM-1 polymers can be found in J. Wu et. a I Adv. Sustainable Syst., 2018, 2, 1800044, of which is herein incorporated by reference. The as-synthesized PIM-1 polymer was functionalized with amidoxime to yield amidoxime- functionalized PIM1 (AOPIM1) by a simple one-pot reaction identically adopted from FI. A. Patel et. a I ., Chem. Commun., 2012, 48, 9989-9991 and is herein incorporated by reference and the ethanol-washed AOPIM1 polymers were dried at 110 °C under vacuum for 20 hours. The respective yields for PIM-1 and AOPIM1 were about 78% and 94%.

Pristine AOPIM1 membrane fabrication

PIM-1 membranes were fabricated by a common solution casting method identical to that in J. Wu et. a I . , Adv. Sustainable Syst., 2018, 2, 1800044. AOPIM1 membranes were also prepared via a simple solution-casting method, followed by a solvent exchange process with MeOFI to remove trapped solvents. Briefly, 0.2 g AOPIM1 polymer was first dissolved in 10 g DMF to prepare a 2 wt% AOPIM1-DMF polymer solution. After stirring overnight and filtering twice using 5 mmPTFE syringe filters, the polymer solution was poured into a Petri dish and then placed in a vented oven at 100 °C and atmospheric pressure for 48 hours to slowly evaporate the solvent. After the film was formed, the Petri dish was removed from the oven and cooled down to room temperature. Deionized water was then added to the Petri dish to immerse the film for 5 min to help delaminate it from the glass surface. The as-cast AOPIM1 film was then cut into even pieces, weighing 65.0 ± 0.5 mg each, and submerged in MeOFI contained in a screw-cap bottle with 100 rpm stirring for 24 hours to drive out the occluded DMF. After solvent exchange with MeOFI, the moisture on the surface of the films was gently wiped off using tissue paper and then the films were dried under vacuum at 120 °C overnight to obtain the pristine AOPIM1 membrane samples. The average sample thickness was 25 ± 5 mm, and the average sample weight loss due to the removal of occluded DMF was 7.8 ± 0.4 mg. One batch of the as-cast AOPIM1 membranes was directly dried under vacuum at 150 °C overnight which was an adequate temperature for the complete removal of trapped DMF solvents without immersion into MeOFI. They were used as control samples to confirm the effective removal of occluded DMF by solvent exchange.

AOPIM1-SCA4 membrane fabrication bv PFI

In screw-cap glass bottles, SCA4 powders, with varied masses of 1.2, 3.1, 5.0 and 10.3 mg, were dissolved into a fixed volume of MeOFI to prepare SCA4-in-MeOFI solutions with different SCA4 concentrations. Then, the as-cast AOPIM1 films of 65.0 ± 0.5 mg each were submerged in these SCA4-in-MeOFI solutions for 24 hours to infiltrate the membranes with SCA4 molecules while exchanging the occluded DMF solvent at the same time. Lastly, these SCA4-infiltrated membrane samples were sent for final drying under vacuum at 120 °C overnight, and the resultant membranes were abbreviated as AOPIMl-SCA4membranes.

Quantification of the degree of infiltration bv titration

The ideal mole ratio of SCA4 to AOPIM1 polymer is defined as the mole ratio of the total SCA4 dissolved in MeOFI to the AOPIM1 polymer content in the as-cast films. Given the average 7.8 mg loss of DMF from 65 mg of the as-cast AOPIM1 films as determined previously, one could then estimate the ideal mole ratios of SCA4 to AOPIM1 polymer to be about 1.5%, 3.8%, 6.2% and 12.7%, respectively. Flowever, not all the SCA4 dissolved in MeOFI could be infiltrated into the membrane samples due to the eventual mass transfer equilibrium reached. Thus, the actual mole ratio of SCA4 to AOPIM1 polymer is defined as the mole ratio of SCA4 in the resultant AOPIM1-SCA4 membranes to the AOPIM1 polymer content in the resultant AOPIM1-SCA4 membranes, and it was measured by the titration of residual SCA4 in the used SCA4-in-MeOFI solutions.

By using a digital pH meter (Oakton p H 450 Meter Kit), the p H values of initial SCA4- in-MeOFI solutions with known SCA4 amounts were measured. Then, these solutions were titrated with a 100 ml MeOFI solution containing 20 mg NaOFI (the titrant) to the point where their pFH values returned to that of pure MeOFI. From the titration results, a calibration curve was established between the SCA4 mass in a fixed volume of MeOFI and the titrant volume used. By interpolating within the calibration curve, the residual mass of SCA4 in the SCA4-in-MeOH solutions used was determined using the same titrant. As such, the actual mass of SCA4 infiltrated into the membranes could be calculated from the mass difference of SCA4 in MeOH before and after the infiltration process. With the dry mass of the resultant AOPIM1-SCA4 membrane samples being measured, the actual mole ratio of SCA4 to AOPIM1 polymer (526 g mol ¹) could be calculated.

The table below summarizes the results. According to the actual mole ratios of SCA4 to AOPIM1 polymer, which were found to be 0.99%, 2.44%, 3.48% and 4.96%, the resultant AOPIM1-SCA4 membranes were named AOPIMl-SCA4-l%, AOPIM1-SCA4- 2%, AOPIMl-SCA4-3% and AOPIMl-SCA4-5%, respectively, for simplicity. These percentage values were also referred to as the 'degree of infiltration' or 'x% infiltration' in this study.

Positron annihilation lifetime spectroscopy (PALS^')

Positron annihilation lifetime spectroscopy (PALS) was used to measure the fractional free volume (FFV) and the average pore radii of the AOPIM1-SCA4 membranes by using a conventional bulk PALS instrument under ambient conditions in the laboratory. The emission (birth) and the subsequent annihilation (death) of a positron from the ²²Na radioactive isotope source were both accompanied by the generation of g-rays (1.28 MeV and 0.511 MeV, respectively) due to the nuclear decay. As such, the detection of the time difference between the birth and death of the positron allowed for the measurement of the lifetime. The setup included two stacks of polymer film samples (each about 0.5 mm thick and l x l cm² in area) sandwiched in between the ²²Na source, and the PALS experiments were performed at a counting rate of 210 to 230 counts per s with a total of 5 x 10^s counts collected for each spectrum. Four fitted lifetime components, namely, TI (para-Positronium (p-Ps)), T2 (free positrons), T3 and T4 (both ortho-Positronium (o-Ps)), were obtained with a small variance from the PALS spectra using the PATFIT program (which assumes a Gaussian distribution of the logarithm of the lifetimes), and the lifetime distribution of these positions was obtained using the MELT program. By approximating the free volume elements as spherical cavities, the mean free-volume radius R A) was calculated from o-Ps lifetimes based on an established semi-empirical correlation as shown below,

where Ti is the o-Ps lifetime, T3 and T4, in nanoseconds (ns) and AR is an empirical constant (1.656 Å). The fractional free volume (FFV) was correlated with R and L· according to the Williams-Landel-Ferry equation,

where Ii = I3 or I4 is the intensity (%) of o-Ps and Ri = R3 or R4 refers to the mean free- volume radius (Å).

Pure- and mixed-aas permeation Pure-gas permeation tests for each membrane sample were conducted in a variable- pressure constant-volume gas permeation cell with gases being run in the order of H₂, O2, N2, and CFI4 to CO2. The detailed setup and procedures can be found in L. Shao et. al., J. Membr. Sci., 2005, 256, 46-56, of which is herein incorporated by reference. After the sample was mounted into the cell and the lid was properly secured, the system was vacuumed at 35 °C for 20 hours before running the first gas. The testing temperature was maintained at 35 °C throughout all the runs.

Eqn (3) shown below, which is based on the steady-state pressure increment (dp/dt), expresses the gas permeability in terms of several measurable parameters,

where P is the gas permeability of the membrane in Barrer (1 Barrer = 1 x 10^-10 cm³ (STP) cm cm^-2 s^-1 cmPig^-1), V represents the downstream reservoir volume (cm³), A denotes the effective membrane area (cm²), I represents the membrane thickness (cm), T represents the operating temperature (K) and lastly på is the upstream pressure (psia). Each gas was tested at least two times to ensure that the deviation is within 1%.

The ideal selectivity of any gas pair is the ratio of the pure gas permeability of one gas to that of the other and was calculated using the following equation,

where OA/B is the ideal selectivity of the membrane for gas species A over B and PA and PB represent the single gas permeability of gas species A and B, respectively.

According to the solution-diffusion model, the permeability of polymer membranes for a particular gas is defined as the product of its diffusivity (D) and solubility (S) for that gas. By manipulating eqn (4), the ideal permeability selectivity, Op, can then be defined as the product of diffusivity selectivity, OD, and solubility selectivity, as, which is expressed by eqn (5) here,

where DA and DB are the diffusivity coefficients (cm² s^-1) and SA and SB are the solubility coefficients (cm³ (STP) per cm³ membrane bar) of gas species A and B, respectively.

A binary gas mixture of 50% PI 2 and 50% CPU (equimolar) was used for mixed gas permeation tests. The operating temperature was again maintained at 35 °C with a continuously supplied feed at 7 bar. The cell was also slowly vented at the upstream side to maintain constant gas composition there. The composition of the permeate was analyzed using an Agilent 7890 gas chromatography (GC) system. The mixed gas permeability of H2 and CPU was determined using the following equations,

where P_CO2 and PCH4 denote the permeability of CO2 and CPU, respectively, P2 is the upstream feed pressure (psia) and x and y represent the mole fractions of the gases in the feed and permeate, respectively. Other symbols have the same meanings as previously described, and the mixed-gas selectivity was calculated as the ratio of their permeability.

Gas sorption measurements

The gas sorption properties of AOPIM1-SCA4 membranes were investigated using an XEMIS-series static sorption microbalance system (UK) using a gravimetric method. Each membrane sample, weighing about 30 to 35 mg, was loaded into the microbalance chamber and the system was first stabilized under < 10^-6 vacuum for 12 hours at 35 °C (maintained using a thermally controlled water bath) after each gas being admitted into the system. The pressure was then gradually increased from 50 mbar to a maximum of 10 bar followed by slow desorption, and the concentration of adsorbed gas (C, cm3 (STP) per cm³ membrane) as a function of the system pressure was computed using the built-in Hisorp software. The three tested gases were run in the order of N2 and CPU to CO2 for each sample and all isotherm data points were obtained with a standard deviation of maximally ±10%.

The solubility coefficient (S, cm³ (STP) per cm³ membrane bar) was calculated according to eqn (8) below,

where p denotes the feed pressure (bar).

The obtained sorption isotherms of AOPIM1 and AOPIM1-SCA4 membranes were fitted to the dual-mode sorption model for glassy polymer as expressed in the equation below,

where C is the gas concentration (cm³ (STP) per cm³ membrane), kD represents the Henry's law coefficient (cm³ (STP) per cm³ membrane bar), and C'H and b are the Langmuir capacity parameter (cm³ (STP) per cm³ membrane) and affinity parameter (bar¹), respectively.

Organic solvent forward osmosis (OSFO) application Preparation of the cross-linked membrane substrate

The fabrication and crosslinking procedures are as follows. Briefly, Matrimid was first dried in a vacuum oven at 70 °C for 24 h before it was dissolved in a mixture of N- methyl-2-pyrrolidinone (NMP) and polyethylene glycol 400 (PEG 400) with a weight ratio of 20:64: 16 for Matrimid, NMP, and PEG400. The mixture was stirred at 70 °C for another 24 h to prepare a homogeneous dope and then degassed for at least 1 day prior to casting the membrane substrate on a glass plate, followed by phase inversion in a deionized (DI) water bath. The as-cast substrate was stored in a DI water bath for at least 1 day to complete the phase inversion. To cross-link the substrate, it was cut into a proper size and immersed in an isopropanol (IPA)/water (50/50 wt/wt) solution containing 5 wt% 1,6-hexanediamine (HDA) for 24 h. Afterward, the cross-linked substrate was taken out, washed thoroughly with fresh DI water and stored in DI water for further modifications.

Fabrication of TFN membranes via interfacial polymerization

Interfacial polymerization between MPD (aqueous phase) and TMC (organic phase) was conducted to form a thin polyamide layer on top of the cross-linked Matrimid membrane substrate, as depicted in Figure 7c. Briefly, the cross-linked membrane substrate was first immersed into a 2 wt% MPD aqueous solution containing 0.2 wt % SDS for 2 min. For each TFN membrane containing STCAss or SCA, a predetermined amount of calixarene was added into the MPD solution. Afterward, the membrane was taken out and the extra MPD solution on the surface was wiped with filter papers. A hexane solution of 0.1 wt% TMC was then deposited onto the top of the membrane for 1 min, followed by air-drying for 5 min to form the dense polyamide layer. The prepared TFC and TFN membranes were then rinsed with ethanol and immersed in ethanol prior to further testing. The prepared membranes were denoted as TFC-0 (pristine), TFN- STCAss-0.5, TFN-STCAss-1, TFNSTCAss-1.5, TFN-STCAss-2, and TFN-SCA-1.5, where 0.5, 1, 1.5, and 2 refer to the weight percentages of STCAss or SCA in the MPD solutions. To prepare free-standing polyamide films, an MPD aqueous solution of ~10 ml with the corresponding compositions and a TMC/hexane solution of ~15 ml were prepared. Firstly, the MPD aqueous solution was poured into a petri dish and allowed to stabilize the liquid surface. Subsequently, the TMC/hexane solution was added dropwise on the top surface of the MPD solution. The petri dish was then covered with a lid to prevent the hexane evaporation and stabilize the film growth. After 24 h, the petri dish was drained and the resultant thin film was rinsed several times with ethanol to remove the excess monomers. It was then vacuum dried for further characterizations.

Solvent reclamation through OSFO

The solvent reclamation was conducted using a lab-scale OSFO unit with solvent resistant tubing and pumps. Both draw and feed solutions were circulated counter- currently using the pumps at volumetric flows of 0.2 L min^-1. Two operating modes; namely, FO mode (i.e., the selective layer facing the feed solution) and PRO mode (i.e., the selective layer facing the draw solution), were studied. The solvent flux (Jw, LMFI) and reverse solute flux (Js, gMFI) were determined using Eqs. (10) and (11), respectively.

Where Am (g) is the absolute weight loss in the feed side or the absolute weight gain in the draw side, p (g cm^-3) is the solvent density, At (h) is the testing duration of 2 h, Am (cm²) is the effective contact area of 4 cm², ΔCt (g L^-1) is the change of solute concentration in the feed solution, and V (L) is the volume of the feed solution.

In OSFO tests, LiCI was employed as the draw solute and was dissolved in pure ethanol at a concentration of 2M. Meanwhile, the pure ethanol was used as the feed solution. The weight gain of the feed solution was monitored using a balance (A&D Company, Ltd, Japan) connected to a data log on a computer. The reverse solute flux was determined using a conductivity meter (Metrohm, Switzerland), where the calibration curve was attained prior to the tests. To demonstrate the feasibility of concentrating pharmaceuticals and recovering organic solvents using the developed TFN membranes in OSFO processes, tetracycline and paracetamol/ethanol solutions of 2000 ppm were used as the model feed solutions. The FO mode was chosen as it had a lower reverse solute flux and a less fouling tendency than the PRO mode. The membrane rejection to tetracycline/ paracetamol, Rf (%), was defined as follows:

Where Cd (g L^-1) is the tetracycline/paracetamol concentration in the draw solution at the end of each OSFO test, Vd (mL) is the final volume of the draw solution, Vp (mL) is the volume of the permeate, and Cf (g L^-1) is the tetracycline/paracetamol concentration in the feed solution. The tetracycline and paracetamol concentrations were determined by a UV-Vis spectrophotometer (Pharo 300, Merck) according to the Beer-Lambert law.

Transport properties The pure solvent permeance (A, Lm² h^_1 bar^-1, LMFI/bar) and salt rejection (Rs, %) were determined by testing the membranes under a transmembrane pressure (DR) of 10 bar in dead-end cells at room temperature. The feed solutions were made of 200 ppm LiCI, NaCI and KCI in ethanol. The concentrations of salt in the feed (Cf, g L^_1) and permeate (Cp, g L^_1) were determined using a conductivity meter (Metrohm, Switzerland). The pure solvent permeance and solute rejection R_s were calculated by Eqs. (13) and (14), respectively.

Where ΔV (L) is the permeate volume, At (h) is the testing duration, A_m is the effective area of the membrane, and DR is the applied transmembrane pressure.

Membranes were also tested using water, methanol, and ethanol under a DR of 10 bar in dead-end cells at room temperature to further understand the solvent transport mechanism inside the pristine and functionalized polyamide layers. The permeability (P, cm² s^_1 bar^-1) of membranes was calculated using Eq. (15)

P = A X Δx ( 15)

Where Dc (m) is the membrane thickness. As the transport resistance of the developed membrane was mainly determined by the polyamide selective layer, the thickness of the polyamide layer was applied here. The solubility of a membrane was measured using the solvent evaporation method. Briefly, once the free-standing polyamide thin film was fabricated, it was dried in a vacuum oven to remove the moisture. The film was quickly weighed (m₀) and immersed in an excessive solvent for 1 week to ensure it being fully saturated. After that, the solvent was allowed to evaporate at room temperature and the weight change was recorded as a function of time. Generally, the weight profile would display three weight loss rates, which indicated the solvent evaporation from (1) the membrane surface, (2) inside the membrane, and (3) the cease of evaporation. Therefore, the difference between the turning point (m_w) of the first and second slopes with the final weight (rrid) is considered as the solvent uptake by the membrane. The solubility of the polyamide thin film (S, gram solvent per gram membrane, g g^_1) was calculated using Eq. (16).

The classical solution-diffusion model works for the case where flux is linearly dependent on transmembrane pressures, and thus the diffusivities (D, cm² s^_1) of the selected solvents could be determined as follows:

Where C_s (gm^-3) is the solvent concentration inside the membrane and V_s (m³ mol^-1) is the partial molar volume of the respective solvent.

Hollow Fiber Membrane application

Fabrication of the Torlon^® outer selective HFM substrate

The Torlon^® polymer and Li Br were vacuum dried overnight at 70 °C prior to preparing the dope solution. LiBr was firstly dissolved in N-methyl-2-pyrrolidinone under stirring at 80 °C, then the Torlon^® polymer was added gradually. The mixture was stirred vigorously until the polymer was dissolved. Once the dope became homogenous and cooled down, tetrahydrofuran was added as the co-solvent. The final dope mixture was degassed prior to being loaded into an ISCO syringe pump before spinning. The hollow fiber substrates were fabricated via dry-jet wet-spinning and the table below summarizes the spinning parameters. The as-spun FIFMs were then immersed in tap water for 2 days to remove residual solvents and to complete the phase inversion process. Prior to making membrane modules, the as-spun FIFMs were immersed into a 50/50 wt% glycerol/water solution followed by air-drying to prevent the pores from collapsing. Each membrane module contained 3 pieces of fibers, each fiber had an effective length of 15 cm.

Crosslinking modifications and SCA8 impregnation

The as-spun FIFMs preserved in water were used for crosslinking modification. A 5 wt% 1,6-hexanediamine (FIDA) or tris(2-aminoethyl)amine (TAEA) solution in 50/50 wt% isopropanol/DI water was prepared prior to use. The membranes were immersed into the solution for 24 h at room temperature. The HDA- or TAEA-crosslinked HFMs were subsequently washed and stored in DI water overnight to remove any excess crosslinker. The HDA- or TAEA- crosslinked HFMs were denoted as HDA-XIinked HFM or TAEA-XIinked HFM, respectively.

The TAEA-XIinked HFMs were further impregnated with SCA8 by means of a solvent infiltration method. Briefly, the TAEA-XIinked HFMs were immersed in a 50/50 wt% methanol/DI water solution containing a certain amount of SCA8 for 1 h. The SCA8 concentration was varied to optimize the OSN performance and the prepared HFM was denoted as TAEA-SCA8-X, where x stands for the SCA8 concentration in the solution. The actual SCA8 loading in the HFMs was determined by measuring the pH difference of the solutions before and after the impregnation. The actual SCA8 loadings in HFMs were calculated to be 0.5 wt%, 1.7 wt% and 2.8 wt% when using SCA8 solutions of 0.1 wt%, 0.3 wt% and 0.5 wt%, respectively.

In detail, the SCA8 concentration in the solution was varied to optimize the OSN performance. To determine the actual SCA8 loading in HFMs, the pH values of the solutions were measured before and after the immersion of HFMs. By calculating the amount of H⁺ inside the solvent, the difference in SCA8 amount in the solutions before and after the immersion could be back-calculated. As a result, the actual SCA8 loadings in HFMs were calculated to be 0.5 wt%, 1.7 wt% and 2.8 wt% when using SCA8 solutions of 0.1 wt%, 0.3 wt% and 0.5 wt%, respectively. A sample calculation of TAEA- SCA8-0.1 is given below.

Since the 0.1 wt% SCA8 in methanol/water solution had an initial pH of 2.728 and a final pH of 2.76 and the solution had a volume of 28 ml_, the amount of H⁺ ions absorbed by the HFMs was about 3.7 μmol (= (io^-2·⁷²⁸ - io^-2·⁷⁶) x 0.028). Because each SCA8 could dissociate 8 H⁺, the SCA8 amount attached onto the HFMs was 0.46 pmol (= 3.7/8). The SCA8 had a molecular weight of 1489.45 g mol ¹ and 6 pieces of HFMs had a weight of 0.14 g, one could calculate the actual SCA8 loading to be 0.5 wt%

Pure water permeance and pore size distribution

The pure water permeance (PWP, L m² h^_1 bar¹, LMH/bar) of the as-spun, HDA-XIinked and TAEA-XIinked HFMs were measured using a cross flow setup at a constant flow rate of 0.5 L min ¹ under a trans-membrane pressure (DR) of 2 bar. The PWP was calculated using the following equation:

where Q (L/h) is the water flow rate at the permeate side and A (m²) is the effective membrane area.

The pore size distribution was determined using the solute rejection method with PEGs. Water solutions containing 200 ppm PEGs were used as the feed solutions. The solution concentrations in both feed ( c_f ) and permeate (c_p) were determined using a total organic carbon analyzer (TOC ASI-5000A, Shimadzu). The solute rejection ( R , %) for each organic solute was calculated using Equation (19) and the Stoke diameters (d_s) of the PEG solutes were determined by Equation (20).

Subsequently, the solute rejection was plotted against Stoke diameter on a log-normal probability graph and the linear regression was performed. The molecular weight cutoff (MWCO) is the molecular weight of a solute where its rejection is 90%, and the mean effective pore diameter ( μ_r ) is the size of a solute where its rejection is 50%. The geometric standard deviation (σ_r) of a membrane is the size ratio of solutes with rejections of 84.13% and 50%. The pore size distribution of the membrane was generated using Equation (21).

where d_p is the effective pore diameter. OSN performance

The OSN performances of the prepared HFMs were tested using a solvent-resistant cross-flow setup at a solvent flow rate of 0.5 L min ¹ under 2 bar. Various organic solvents were circulated at the shell side of the membrane and the pure solvent permeance was determined using Equation (18). The rejection was obtained using RB, VBB, MB and paracetamol in methanol at a concentration of 50 ppm. The solute rejections towards RB, VBB, MB and paracetamol were calculated using Equation (10). The concentrations of both feed and permeate solutions were determined using a UV- Vis spectrophotometer (Pharo 300, Merck) according to the Beer-Lambert law. For determination of permeance and rejection in either aqueous or organic solutions, the membranes were stabilized for 1 h before taking any measurements. To conduct the 7- day stability test, TAEA-SCA8-0.3 was chosen as the representative. The FIFMs were immersed in MB/methanol for 7 days and the permeance and rejection were collected every day. For all the performance tests, the average values of two permeate samples from each membrane with at least three membrane samples were reported.

The feed solution for the mixed solutes separation is prepared by mixing a 50 ppm DMNA/methanol and a 50 ppm MB/methanol solution together. The UV spectra of the feed and permeant were measured and recorded. The separation factor is determined using the following equation.

Fabrication of the PBI FIFM substrate with an outer selective layer A polymer solution of PBI/DMAc/LiCI/Propanol/PVP 10K (18.5/59.5/1.0/18.0/3.0 wt%) was prepared by diluting the original PBI/DMAc dope with additional DMAc (solvent), propanol (non-solvent) and PVP 10K (pore former). The mixture was stirred overnight under 50 °C to obtain a homogenous solution. The purpose of adding propanol into the dope was to facilitate solvent evaporation from the nascent hollow fiber surface in the air-gap region and accelerate the growth of a thin selective outer layer due to its high volatility. Afterwards, the PBI solution was degassed and loaded into a 500 mL ISCO syringe pump, it was further degassed overnight before spinning. A mixture of DMAc/IPA (14/86 wt%) was used as the bore fluid because it had a closer solubility parameter with the spinning dope. It could not only lower the phase separation process in the lumen side but also enhance the sublayer porosity. The PBI solution and bore fluid were precisely metered and pumped into the outer and the inner annuluses of the spinneret, respectively. They travelled an air gap distance of 2.5 cm and then entered an IPA coagulation bath at a temperature of 10 °C. The solidified PBI HFMs were taken up by a rotational drum, cut and washed in tap water for 48 h to leach out residual solvents. Subsequently, a solvent exchange with IPA was conducted. The resultant PBI HFMs were kept in IPA for characterizations and post-modifications. The table below lists the detailed hollow fiber spinning conditions.

DBX crosslinkina modification and SCA8 impregnation

The as-spun PBI HFMs were chemically crosslinked by immersing them in a 5.0 wt% DBX solution in acetonitrile (MeCN) at 80°C for 24 h with continuous stirring and refluxing. The resultant PBI-DBX HFMs were rinsed with MeCN to remove remaining reagents and then were preserved with PEG200 by immersion in a PEG200/IPA (1:1) solution for 24 h to maintain the pore structure prior to air-drying for storage. Some DBX crosslinked PBI HFMs were modified by SCA8 further via a solvent infiltration method. Briefly, the PBI-DBX HFMs were dipped in a 50/50 wt% methanol/DI water solution consisting of 0.3wt% SCA8 for 4 h. Then, the SCA8 modified HFMs were rinsed by IPA, followed by immersing them in a PEG200/IPA (1:1) solution for 24 h and airdrying for further use. The resultant membrane is referred to as PBI-DBX-SCA8 HFMs.

Fabrication of 1-inch HFM modules and long-term tests

The PBI-DBX-SCA8 hollow fiber bundle was placed in a stainless steel tubing assembled with Swagelok fittings and sealed with removable potting caps. The assembled module was horizontally mounted on a centrifugal potting instrument. The epoxy potting solution was radially pushed to cram the mold under the centrifugal force and formed the tubesheet until being solidified. The fabricated 1-inch PBI HFM module had a packing density of 52% and an effective outer area of around 0.22 m². A feed of 10 wt% oil in acetone was used to evaluate the long-term stability on a crossflow setup at a flow rate of 1.5 L min ¹ under 10 bar for 30 days. The permeant was recirculated into the feed container to keep the feed concentration constant. The oil concentrations of the permeant and feed were measured and recorded.

It will be appreciated that many further modifications and permutations of various aspects of the described embodiments are possible. Accordingly, the described aspects are intended to embrace all such alterations, modifications, and variations that fall within the spirit and scope of the appended claims.

Throughout this specification and the claims which follow, unless the context requires otherwise, the word "comprise", and variations such as "comprises" and "comprising", will be understood to imply the inclusion of a stated integer or step or group of integers or steps but not the exclusion of any other integer or step or group of integers or steps.

Throughout this specification and the claims which follow, unless the context requires otherwise, the phrase "consisting essentially of", and variations such as "consists essentially of" will be understood to indicate that the recited element(s) is/are essential i.e. necessary elements of the invention. The phrase allows for the presence of other non-recited elements which do not materially affect the characteristics of the invention but excludes additional unspecified elements which would affect the basic and novel characteristics of the method defined. The reference in this specification to any prior publication (or information derived from it), or to any matter which is known, is not, and should not be taken as an acknowledgment or admission or any form of suggestion that that prior publication (or information derived from it) or known matter forms part of the common general knowledge in the field of endeavour to which this specification relates.

Claims

1. A membrane composite, comprising : a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; and wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer.

2. The membrane composite according to claim 1, wherein the first polar moiety resides in a polymer backbone, as a side group, or a combination thereof.

3. The membrane composite according to claim 1 or 2, wherein the first polar moiety is selected from cyano, acyl, oxyacyl, acyloxy, amino, acylamino, aminoacyl, amidoximyl, oximyl, hydrazonyl, iminyl, keto oximyl, aceto oximyl, hydroxyl, pyrrolyl, oxazolyl, imidazolyl, pyrazolyl, triazolyl, and a combination thereof.

4. The membrane composite according to any one of claims 1 to 3, wherein the macrocycle or derivative thereof is selected from cyclodextrin, calixarene, cucurbituril, resorcinarene, pillararene, and a combination thereof.

5. The membrane composite according to any one of claims 1 to 4, wherein the macrocycle or derivative thereof is a calixarene or a derivative thereof, selected from sulfocalixarene (or sulfonylcalixarene), sulfothiacalixarene, carboxylatocalixarene, aminocalixrene, p-phosphonic acid calixarene, or a combination thereof.

6. The membrane composite according to any one of claims 1 to 5, wherein the macrocycle or derivative thereof has monomer residues of 4 to 12.

7. The membrane composite according to any one of claims 1 to 6, wherein the macrocycle or derivative thereof has an upper rim and a lower rim, the upper rim and lower rim are separated by a frustum-shaped cavity, wherein a diameter of the lower rim is at least about 3 A.

8. The membrane composite according to claim 7, wherein the upper rim is functionalised with at least one moiety independently selected from sulfonyl, phosphoryl, amino, carboxyl, oxyalkyl and a combination thereof.

9. The membrane composite according to claim 7 or 8, wherein the lower rim is functionalised with at least one moiety independently selected from optionally substituted acyloxy, optionally substituted acyl, optionally substituted alkyl, optionally substituted alkenyl, optionally substituted alkoxy, and optionally substituted alkenyloxy.

10. The membrane composite according to any one of claims 1 to 9, wherein a mole ratio of macrocycle to polymer is about 0.1% to about 10%.

11. The membrane composite according to any one of claims 1 to 10, wherein the macrocycle is about 0.1 wt% to about 10 wt% relative to the membrane composite.

12. The membrane composite according to any one of claims 1 to 11, wherein the macrocycle is homogenous dispersed across the thickness of the layer of polymer.

13. The membrane composite according to any one of claims 1 to 12, wherein the macrocycle is further covalently bonded to the layer of polymer.

14. The membrane composite according to any one of claims 1 to 13, wherein the polymer in the layer of polymer is selected from a polymer with intrinsic microporosity- 1 (PIM-1), polyamide, polyamide-imide, polybenzimidazole, polylactic acid (PLA), polycarboxyl ic acid (PCA), polyethylenimine (PEI), polyarylamine, polyalkylamine, polyallylamine, poly(vinyl amine) or a combination thereof.

15. The membrane composite according to any one of claims 1 to 14, wherein the polymer layer of polymer is crosslinked.

16. A method of fabricating a membrane composite, the membrane composite having a macrocycle homogenously distributed within a layer of polymer, the method comprising incubating the layer of polymer in a macrocycle solution; wherein the polymer comprises a first polar moiety; and wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer.

17. The method according to claim 16, wherein the macrocycle is homogenously dispersed across the thickness of the layer of polymer.

18. The method according to claim 16 or 17, wherein the electrostatic interaction is selected from hydrogen-bond, proton transfer, ionic interaction or a combination thereof.

19. The method according to any one of claims 16 to 18, wherein the macrocycle solution comprises a protic solvent.

20. The method according to claim 19, wherein the protic solvent is selected from methanol, ethanol, isopropanol, n-propanol, water, and a combination thereof.

21. A membrane composite, comprising : a) a layer of polymer; and b) a macrocycle homogenously distributed within the polymer layer; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from acyl, amino, acylamino, aminoacyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 5 to 8.

22. The membrane composite according to claim 21, wherein the layer of polymer is crosslinked with a crosslinker having at least 2 pendant groups, the pendant group selected from aminoalkyl, haloalkyl, and a combination thereof.

23. The membrane composite according to claims 21 or 22, wherein the polymer is selected from polyamide-imide or polybenzimidazole.

24. The membrane composite according to any one of claims 21 to 23, wherein the macrocycle is sulfocalix[n]arene.

25. The membrane composite according to any one of claims 21 to 24, wherein the macrocycle is 4-sulfocalix[8]arene (SCA8).

26. The membrane composite according to any one of claims 21 to 25, wherein a weight ratio of macrocycle to polymer is about 0.1 wt% to about 5 wt%.

27. The membrane composite according to any one of claims 21 to 26, wherein the macrocycle is homogenous dispersed across the thickness of the layer of polymer.

28. The membrane composite according to any one of claims 21 to 27, having a thickness of about 80 pm to about 150 pm.

29. The membrane composite according to any one of claims 21 to 28, having a FTIR spectrum peak of about 1040 cm ¹.

30. The membrane composite according to any one of claims 21 to 29, having a sulphur surface composition of at least about 1 At%.

31. The membrane composite according to any one of claims 21 to 30, having a free volume intensity S-parameter of at least about 0.455.

32. The membrane composite according to any one of claims 21 to 31, having a pore size distribution R-parameter of at least about 0.465.

33. The membrane composite according to any one of claims 21 to 32, having a Victoria blue B rejection of at least about 90% when the macrocycle loading is about 0.1 wt%.

34. The membrane composite according to any one of claims 21 to 33, having a methylene blue rejection of at least about 80% when the macrocycle loading is about 0.1 wt%.

35. The membrane composite according to any one of claims 21 to 34, having a paracetamol rejection of at least about 30% when the macrocycle loading is about 0.1 wt% .

36. The membrane composite according to any one of claims 21 to 35, having a pure methanol permeance of about 1 LMH/bar to about 1.8 LMH/bar when the macrocycle loading is about 0.3 wt%.

37. The membrane composite according to any one of claims 21 to 36, having a pure acetonitrile permeance of about 1 LMH/bar to about 0.28 LMH/bar when the macrocycle loading is about 0.3 wt%.

38. The membrane composite according to any one of claims 21 to 37, having a pure acetone permeance of about 0.4 LMH/bar to about 1 LMH/bar when the macrocycle loading is about 0.3 wt%.

39. The membrane composite according to any one of claims 21 to 38, having a pure ethanol permeance of about 0.48 LMH/bar to about 0.76 LMH/bar when the macrocycle loading is about 0.3 wt%.

40. The membrane composite according to any one of claims 21 to 39, having a pure ethyl acetate permeance of about 0.24 LMH/bar to about 0.36 LMH/bar when the macrocycle loading is about 0.3 wt%.

41. The membrane composite according to any one of claims 21 to 40, having a pure tetrahydrofuran permeance of about 0.12 LMH/bar to about 0.24 LMH/bar when the macrocycle loading is about 0.3 wt%.

42. The membrane composite according to any one of claims 21 to 41, having a pure toluene permeance of about 0.01 LMH/bar to about 0.08 LMH/bar when the macrocycle loading is about 0.3 wt%.

43. The membrane composite according to any one of claims 21 to 42, having a stability in a methylene blue/methanol mixture for at least 7 days.

44. The membrane composite according to any one of claims 21 to 43, having a N,N- dimethyl-4-nitroaniline (DMNA)/methylene blue (MB) separation factor of about 14.5.

45. The membrane composite according to any one of claims 21 to 44, being formed as a hollow fiber membrane.

46. A method of fabricating a membrane composite, the membrane composite having a macrocycle homogenously distributed within a layer of polymer, the method comprising incubating the layer of polymer in a macrocycle solution; wherein the polymer comprises a first polar moiety; wherein the macrocycle comprises a second polar moiety in order to form an electrostatic interaction with the first polar moiety of the polymer; wherein the first polar moiety is selected from acyl, amino, acylamino, aminoacyl, and a combination thereof; wherein the macrocycle is calix[n]arene or a derivative thereof; and wherein n is an integer selected from 5 to 8.

47. The method according to claim 46, further comprising a step of crosslinking the layer of polymer with 1,6-hexanediamine (HDA), tris(2-aminoethyl)amine (TAEA), a,a'- dibromo-p-xylene or a combination thereof.

48. The method according to claim 47, wherein the crosslinking is performed by incubating the layer of polymer in a crosslinker solution.

49. The method according to claim 48, wherein the crosslinker solution has a crosslinker concentration of about 5 wt%.

50. The method according to any one of claims 48 or 49, wherein the crosslinking step is performed for at least 12 h.

51. The method according to any one of claims 46 to 50, wherein the macrocycle is homogenous dispersed across the thickness of the layer of polymer.

52. The method according to any one of claims 46 to 51, wherein the macrocycle solution comprises a protic solvent.

53. The method according to claim 52, wherein the protic solvent is selected from methanol, ethanol, water or a combination thereof.

54. The method according to any one of claims 46 to 53, wherein the layer of polymer is incubated in a macrocycle solution for at least 0.5 h.

55. The method according to any one of claims 46 to 54, further comprising a drying step after the incubation step.