Microporous carbon from fullerene impregnated

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Cite this: DOI: 10.1039/c8ta01200h

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Microporous carbon from fullerene impregnated porous aromatic frameworks for improving the desalination performance of thin film composite forward osmosis membranes† Xing Wu,abc Mahdokht Shaibani,cd Stefan J. D. Smith, c Kristina Konstas,c Matthew R. Hill, ce Huanting Wang, e Kaisong Zhang *ab and Zongli Xie

*c

Porous additives and polymer modifications are becoming increasingly common routes to address some of the current challenges faced by membrane technologies. Here, using carbonization and pre-impregnation of fullerene to enhance the performance of a porous aromatic framework, we have developed a novel hydrophobic porous membrane additive (C60@PAF900) for enhancing the performance of forward osmosis (FO) water purification membranes. We examined the influence of C60@PAF900 loading on the polyamide (PA) layer in the fabrication of novel thin film nanocomposite (TFN) FO membranes and subsequently investigated the desalination performance of the resulting TFN FO membranes. The results showed that the addition of nanoporous C60@PAF900 greatly enhanced the water flux and water permeability of TFN FO membranes. Using 2 mol L1 NaCl as the draw solution and 10 mmol L1 NaCl as the feed solution, adding 0.01 w/v% C60@PAF900 into the FO membrane increased the water flux from 7.4 LMH to 12.4 LMH in the ALFS mode (active layer facing the feed solution) and from 12.6 LMH to 21.3 LMH in the ALDS mode (active layer facing the draw solution). To understand these results, the Received 5th February 2018 Accepted 14th May 2018

effect of C60@PAF900 on the surface morphology and chemistry of TFN FO was also investigated. Compared to the pristine FO membranes, the TFN FO membranes possessed a less dense crosslinked network due to the influence of C60@PAF900 on the interfacial polymerization process, and the

DOI: 10.1039/c8ta01200h

interfacial repulsion between C60@PAF900 and polyamide. Together, this study provides an insight into

rsc.li/materials-a

the performance and potential effects of C60@PAF900 on advanced TFN FO membranes.

Introduction Global water scarcity is one of the most severe crises faced by mankind today.1 Thin lm composite (TFC) membranes are commonly used in nanoltration (NF), reverse osmosis (RO) and forward osmosis (FO) technologies for wastewater management and seawater desalination.2–4 FO-based desalination systems have lower energy cost5 and higher water recovery than RO and NF ltration, though some challenges remain.6 The water ux and salt rejection of existing TFC FO membranes

a

Key Laboratory of Urban Pollutant Conversion, Institute of Urban Environment, Chinese Academy of Sciences, Xiamen, 361021, China. E-mail: [email protected]

b

University of Chinese Academy of Sciences, Beijing, 100049, China

c

CSIRO Manufacturing, Private Bag 10, Clayton South, Victoria 3169, Australia. E-mail: [email protected] d

Nanoscale Science and Engineering Laboratory (NSEL), Department of Mechanical and Aerospace Engineering, Monash University, Clayton, VIC 3168, Australia

e

Department of Chemical Engineering, Monash University Clayton, VIC 3168, Australia

† Electronic supplementary information (ESI) available: Raman spectrum of C60@ PAF900; XRD spectrum of C60@PAF900; FTIR spectra of C60@PAF-1 and C60@ PAF900. See DOI: 10.1039/c8ta01200h

This journal is © The Royal Society of Chemistry 2018

require further improvement before the process is efficient. TFC membranes consist of a porous substrate layer coated with a thin polyamide (PA) selective layer, commonly prepared by interfacial polymerization. As the transport of water and salt through the membrane is largely determined by the selective layer performance, research toward improving the PA layer is a promising way to further enhance the efficiency of FO membrane processes. Consequently, many researchers have explored PA layer modication as a possible approach to improving the performance of TFC membranes. Embedding hydrophilic nanoparticles such as silica,7 silver,8 TiO2,9 ZnO,10 MCM-41 NPs,11 and graphene oxides12–14 into the PA layer has been shown to improve the water ux of thin lm nanocomposite (TFN) FO membranes. In these nanocomposites, the additive hydrophilicity increases the solubility of water molecules, which are then able to diffuse through the interfacial void space created by the additive's disruption of the PA interfacial polymerization.15–17 Similar effects can be achieved by the inclusion of porous nanomaterials, which introduce additional void spaces (nanoparticle pore volume) for penetrant transport.18

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Interestingly, recent simulation and experimental research has revealed that improved water transportation can also be achieved when hydrophobic porous materials are added to TFN membranes, demonstrated with zeolites,15,19 MFI zeolites,20 carbon nanotubes (CNTs),21–23 and some metal–organic frameworks (MOFs).24–26 A major factor behind water ux enhancement in these composite materials is that the hydrophobic walls of pores in these materials do not form hydrogen bonds with water molecules. As a result, water molecules could pass through the pores more quickly, oen without reducing salt rejection.15 To capitalize on this mechanism, many researchers are now using hydrophobic porous materials to improve the water ux of membranes. Duan and co-workers15 added zeolitic imidazolate framework 8 (ZIF-8), a hydrophobic MOF, into the organic phase before the interfacial polymerization. The resulting TFN RO membranes showed higher water ux due to water transportation through the MOF's pores. Corry and coworkers21 investigated the effect CNT (1.1 nm diameter; 8,8) functionalization had on membrane water ux and ion rejection, using hydrophilic functionalized (COO, NH3+, OH and CONH2) CNTs. Water uxes were lower for the functionalized CNTs than the native CNTs due to the formation of hydrogen bonds between functionalized CNTs and water molecules. Unfortunately, in this work CNTs also compromised the membrane's salt rejection compared to functionalized CNTs. Humplik and co-workers20 incorporated two types of MFI zeolites into the active layer of RO membranes. Under FO conditions, they found that the hydrophobic MFI zeolite enhanced water transport more than the hydrophilic MFI zeolite. This result was explained by the higher diffusivity of water molecules through the hydrophobic zeolite's pores. Importantly, each of these studies demonstrate that ux enhancement is not limited to hydrophilic additives and can be signicantly enhanced by the inclusion of porous materials. In the gas storage and membrane separation eld, the porous aromatic framework (PAF-1), an amorphous network comprised of carbons linked with biphenyl linkers in a tetrahedral arrangement,27,28 has attracted considerable interest because of its very large surface area and pore volume.27,29,30 As a membrane additive, PAF-1 has been particularly benecial in glassy polymer membranes for gas separation, where its large open pores led to both enhanced gas transport and reduced physical aging of the glassy polymer.31,32 PAF-1's adsorbent and membrane additive performance has been further improved by incorporating lithiated fullerene into the framework's large pores.30,33 More recently, the absence of any hetero-atoms or metals in the PAF-1's structure was exploited by using PAF-1 for the preparation of supercapacitor electrodes.34,35 This was achieved by the graphitization of PAF-1 using a thermal treatment, resulting in Microporous Onion Like Carbons (MOLC).35 Unlike previous examples of framework-derived carbons,36 carbonized PAF-1 retained its high porosity (BET 1084 m2 g1) and did not ˚ in the require a secondary source of carbon.35 Pores of 5 A ˚ in native PAF-1, exhibited resulting carbon, as compared to 12 A excellent size-selectivity and performance for small electrolyte species.35 The characterization of the carbonized PAF-1 revealed a defective graphitic carbon caused by partial collapse of the

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biphenyl-polymer framework. Given the eminence of graphene and graphene oxide (GO) in membrane applications,37,38 we questioned whether a porous graphitic nanoparticle would also offer signicant performance advantages to TFN membrane performance. The carbonized PAF-1 was highly hydrophobic, porous, and although not useful here, electronically conductive; properties all common with graphene and GO. As such, we hypothesized that a carbonised PAF-1 framework may have be an effective hydrophobic additive for water purication membranes. Here, we investigate the applicability of PAF-1 in membranes for water separation, by using framework carbonisation as a method of producing defective graphitic carbons with subnanometre pore sizes. For this work, we included C60 fullerene into PAF-1 (forming C60@PAF) to help template the formation of MOLC during carbonisation at 900 C. The resulting C60@PAF900 nanoparticles were then incorporated into a TFN membrane by dispersing into the organic phase during the interfacial polymerization of polyamide (PA) thin-lm membranes. C60@PAF900-polyamide TFN membranes were prepared with different concentrations of C60@PAF900 to investigate the effect of surface hydrophilicity and membrane morphology on desalination performance by forward osmosis. Our ndings reveal that the novel C60@PAF900 additive inuences the morphology and to a lesser degree, the chemistry of the PA layer in TFN FO membranes, resulting in enhanced water permeability. Moreover, our work provides a novel approach for the design of advanced TFN FO membranes.

Experimental Materials Polysulfone (PSf, Solvay P3500) was purchased from BASF Co., Ltd. (China). 1-Methyl-2-pyrrolidinone (NMP, $99%) and n-hexane (99%) were obtained from Sinopharm Chemical Reagent Co., Ltd. (China). Trimesoyl chloride (TMC, 98%), m-phenylenediamine (MPD) and polyvinylpyrrolidone (PVP K30) were purchased from Sigma-Aldrich (China). 1,5-Cyclooctadiene, bis(1,5-cyclooctadiene)nickel, 2,2-bipyridyl, and tetrakis(4-bromophenyl)methane were purchased from SigmaAldrich (Australia). Sodium chloride (NaCl, Merck, Germany) was used for membrane performance testing.

Preparation of C60@PAF900 Briey, PAF-1 was prepared by Yamamoto coupling of tetrakis(4bromophenyl)methane according to the methodology reported by Zhu and co-workers,29 and as reported in our previous studies.33,35,39 Fullerene (C60) was then impregnated into PAF-1 according to a literature procedure to obtain C60@PAF-1, an offwhite powder with a BET surface area of z2350 m2 g1.30,32 Single step carbonization of C60@PAF-1 was then performed at 900 C for 3 h under nitrogen (120 ml min1) using a heating rate of 2 C min1 (Fig. 1). Aer natural cooling, the resulting carbonaceous material, denoted as C60@PAF900, was characterized and used without any further treatment.


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C60@PAF900 were collected by using a Phillip 1140 diffractometer with a Cu-Ka source (40 kV, 25 mA) at a scan rate of 1 min1 and a step size of 0.02 . Additionally, a Renishaw confocal micro-Raman spectrometer equipped with a HeNe (632.8 nm) laser was applied to investigate the Raman spectra of C60@PAF900. The morphology of C60@PAF900 was investigated by transmission electron microscopy (TEM, FEI Tecnai F20). The surface area and pore size distribution of C60@PAF900 were determined from the N2 adsorption isotherms at 77 K collected using a Micromeritics ASAP 2420 aer degassing samples overnight at 150 C. Fig. 1

Synthesis of C60@PAF900 from C60 impregnated PAF-1.


Characterization of TFC and TFN membranes Preparation of TFC and TFN membranes The PSf substrate membrane was prepared through a typical phase inversion technique as reported in previous studies.13 In brief, 0.5 wt% PVP and 17.5 wt% PSf were dissolved in 82 wt% NMP. The casting solution was stirred with a magnetic stirrer at 60 C. Aer dissolution, the resulting casting solution was cast on a clean glass plate at a casting knife height of 175 mm at ambient temperature. Aer casting, the polymer coated glass plate was immediately immersed into a DI water bath to induce phase inversion. The resulting PSf membrane was then peeled off and immersed into another DI water bath for 24 h before use. The polyamide (PA) layer of FO membranes was prepared by interfacial polymerization directly onto the PSf membrane surface. First, the substrate membrane was immersed in a 2.0 w/v% MPD aqueous solution. Aer 2 min, the MPD aqueous solution was poured off, and excess MPD was removed by using a rubber roller. Aer that, 0.1 w/v% TMC/n-hexane solutions containing different concentrations of C60@PAF900 (0, 0.005, 0.01, 0.015 and 0.02 w/v%) were poured onto the surface of the membrane. Aer 1 min of reaction, the TMC solution was poured off with the resulting composite membrane fully cured by placing in an oven at 60 C for 8 min. Aerwards, the prepared TFC/TFN membranes were stored in a deionized (DI) water bath for future use. The membranes used in this study were named according to the concentration of C60@PAF900 in TMC solution, specically M0 (0 w/v%), M1 (0.005 w/v%), M2 (0.01 w/v%), M3 (0.015 w/v%) and M4 (0.02 w/v%). In order to measure the real amount of C60@PAF900 in TFN FO membranes, the percentage of carbon in TFN FO membranes was determined by using an Elementar elemental analyzer (VARIO MAX), and then compared to that in the TFC FO membrane. The results in Table S1† indicated that the weight ratio between C60@PAF900 and the membrane for M1–M4 membranes was 0.33, 0.52, 0.78 and 1.47 wt%, respectively.

Characterization of C60@PAF900 The carbonization of C60@PAF900 was investigated using attenuated total reectance Fourier transform infrared spectroscopy (ATR-FTIR, Perkin-Elmer, USA) with a range of 500– 4000 cm1. Powder X-ray diffraction (PXRD) patterns of


The characterization of functional groups on the PA surface of the prepared FO membranes was performed using an ATR-FTIR (Nicolet 6700 FTIR, Thermo Fisher Scientic Inc., USA), aer drying the stored membrane samples at room temperature. The hydrophilicity of the polyamide membrane surface was investigated through the water contact angles using a sessile drop analysis system (CAM200, KSV, Finland). For each membrane, the water contact angle was measured every 5 s, over a period of 300 s. To minimize experimental errors, reported contact angle values are averaged from three randomly selecting locations on each sample. The morphology of the prepared FO membranes was investigated using a eld emission scanning electron microscope (FEI Nova NanoSEM 450) system. Evaluation of the desalination performance of FO membranes The water permeability (A) and the salt permeability (B) coefficients of FO membranes were determined using a bench-scale dead-end ltration test system described in previous studies.40,41 The effective membrane area was 14.6 cm2 and the operating pressure was maintained at 0.2 MPa. DI water (0.5 L) and 20 mmol L1 NaCl solution (0.5 L) were used as the feed solutions to acquire the A and B values, respectively. A and B were calculated using eqn (2) and (4), respectively. Reported values are averaged from three samples from each FO membrane. J¼ A ¼

DV Am Dt

(1)

J DP

(2)

where J is the water ux, DV is the volume of permeated water, Am is the membrane area, Dt is the permeation time, and DP is the ltration pressure difference. Cp R ¼ 1 10 (3) Cf 1R B ¼ R AðDP DpÞ

(4)

where R is the salt rejection, Cp and Cf are the NaCl concentrations on the permeate solution and the feed solution, respectively and Dp is the osmosis pressure difference across the membrane.

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The desalination performance of the prepared membranes was evaluated using a cross-ow forward osmosis testing system (Fig. S1†),42,43 in both the ALFS mode (the active layer facing the feed solution) and the ALDS mode (the active layer facing the draw solution) at room temperature. The total active membrane area was 33.8 cm2. 2 L of 2 mol L1 NaCl and 10 mmol L1 NaCl solutions were used as the draw and feed solution, respectively. Aer running for 20 min, the FO system was stable. The water ux (JW) and the reverse solute ux (JS) were measured based on the weight change of the draw solution and the conductivity change of the feed solution, respectively. Water ux and reverse solute ux values are averaged from three different membrane samples, each tested ve times, with JW and JS values measured every 10 min according to the following equations: DV DtAm

(5)

DðCt Vt Þ DtAm

(6)

JW ¼ JS ¼

where DV is the change in the volume of permeated water over a permeation time Dt, Am is the membrane area, Ct is the concentration of NaCl at the end of permeation time, and Vt is the volume of permeated water at the end of permeation time.

Results and discussion

concentric surfaces in C60@PAF900. The distance between two ˚ which matches the concentric layers was approximately 3.4 A, 35 (002) lattice parameter of bulk graphite. The carbonization of C60 impregnated PAF-1 to defective graphite was conrmed by Raman spectroscopy (Fig. S2†) and X-ray diffraction (Fig. S3†). Two characteristic peaks were observed at 1326 and 1582 cm1 in the Raman spectrum, representing the D band and G band of graphite. Moreover, the XRD pattern of C60@PAF900 was mostly amorphous, except for two peaks typical of graphitic carbon at 2q z 24 and 43.2 .35 The concentric graphitic structure of C60@PAF900 matches that of the previously described microporous onion-like carbon (MOLC) prepared directly from PAF1.35 The visual comparison of TEM images from that study suggests that the addition of C60 into PAF-1 before carbonization may encourage the formation of defective graphitic onions by providing a curved graphitic ‘seed’ for the structural rearrangement of PAF-1's biphenyl struts, as discussed by Shaibani and co-workers.35 The porosity of C60@PAF900 was investigated by measuring its N2 adsorption isotherms at 77 K. A typical type-I isotherm curve was observed (Fig. 3a),35,44 conrming C60@PAF900's microporosity. As previously reported, the pore size distribution showed that the majority of C60@PAF900 pores were centred at ˚ with the remainder at about 12 A, ˚ the main pore around 5 A, size in PAF-1(Fig. 3b). Although much lower than that of PAF-1 or C60@PAF-1, the BET surface area (608.3 m2 g1) of

Characterization of C60@PAF900 and TFN membranes Fig. 2 shows the TEM images of C60@PAF900. The high-magnication TEM image (Fig. 2b) indicated the presence of

Fig. 2

TEM images of C60@PAF900.

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Fig. 3 (a) N2 adsorption isotherm of C60@PAF900 at 77 K, and (b) pore size distribution of PAF-1 and C60@PAF900 calculated using the DFT model.


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Fig. 4 (a) FTIR spectra of TFC and TFN FO membranes and (b) interfacial polymerization reaction of TMC and MPD to form the crosslinked polyamide selective layer. Disruption of the polymerization reaction results in a greater proportion of the linear repeat unit within the polymer structure.

C60@PAF900 conrms that PAF-1's porosity has been retained during carbonization as previously reported.35 Fig. 4a illustrates the FTIR spectra of the prepared FO membranes. Some typical characteristic peaks were observed in the spectra of all FO membranes prepared by the interfacial polymerization of MPD and TMC. Peaks at 1660 cm1 and 1542 cm1 are associated with the C]O stretching of the amide

Fig. 5

I band and C–H stretching of the amide II band, respectively. The peak at 1411 cm1 was due to the C–H symmetric deformation vibration of –C(CH3)2 groups in the polysulfone substrate.45 Compared with the (M0) TFC membrane, the TFN membranes (M1–M4) showed increasing transmittance at 1610 cm1 with C60@PAF900 content (Table S2†). This peak is associated with free –COOH functional groups and is indicative of a less connected amide network,46–48 suggesting that the inclusion of C60@PAF900 disrupts the interfacial polymerization reaction (Fig. 4b). The presence of C60@PAF900 and the morphology of FO membranes were investigated using SEM. All selected FO membranes (M0, M2 and M4) showed the characteristic “ridge and valley” structure of polyamide membranes fabricated by the interfacial polymerization (Fig. 5). In the SEM images of M2 (Fig. 5c and d) and M4 (Fig. 5e and f) membranes, C60@PAF900 can be observed both on the top surface of the PA layer (“ridges”) and on the surface covered with polyamide (“valleys”), highlighted by red and yellow arrows, respectively. This indicates that C60@PAF900 is distributed throughout the semicontinuous PA layer. Compared with the M4 membrane (0.02 w/v% of C60@PAF900), the M2 membrane (0.01 w/v% of C60@PAF900) shows a better dispersion of C60@PAF900 on the membrane surface. High magnication SEM images of the M0's polyamide surface show dense cross-linked networks with few globular structures (Fig. 5b). A similar surface morphology is observed for the M2 membrane. However when the concentration of C60@PAF900 is increased to 0.02 w/v% for M4, the crosslinked networks became sparser, with more globular structures and C60@PAF900 aggregates appearing on the surface of the membranes (Fig. 5f). Cross-sectional SEM images of the prepared FO membranes show that the PA layers in the M2 and M4 membranes were thicker than that in the M0 membranes

Surface SEM images of the (a and b) M0 membrane, (c and d) M2 membrane and (e and f) M4 membrane.


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Cross-sectional SEM images of the (a) M0 membrane, (b) M2 membrane and (c) M4 membrane.

Fig. 6

(Fig. 6). Fig. 6b shows a C60@PAF900 aggregate on the membrane surface using high-magnication SEM. Together these results indicated that the presence of C60@PAF900 and C60@PAF900 aggregates caused an increase in the thickness and linearity of the PA layer in the TFN membranes. These morphological differences between the TFC and TFN FO membranes can be explained by the inuence of C60@PAF900 on the interfacial polymerization reaction. The formation of the crosslinked network was caused by the diffusion of MPD into the organic phase, and the subsequent reaction of MPD with TMC during the interfacial polymerization.13,15 Critically, the extent of

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interfacial polymerization is limited by the diffusion of MPD through the barrier layer formed at the interface of the organic and inorganic phases. Once the interfacial polymerization forms a barrier to MPD, polymerization effectively stops and the thickness of the PA layer of membrane stabilizes.15 When C60@PAF900 was included in the organic phase, gradual settling of the additive causes hydrophobic C60@PAF900 to accumulate on the substrate surface and within the PA layer as it forms (Fig. 7a). The included C60@PAF900 then acts as a barrier to MPD diffusion from the inorganic phase, as shown by the orange arrows in Fig. 7b. This restriction of MPD diffusion to the TMCcontaining organic phase disrupts the polyamide chain propagation and crosslinking within the local phase interfacial region. Interestingly, despite the in situ polymerization reaction, it appears that C60@PAF900 is not in intimate contact with the formed PA layer and instead forms a defective polymer-additive interface. This may be explained by TMC's p–p interaction with C60@PAF900 that orientates the conformation of the forming polymer away from the additive's surface, as observed in other ‘poorly compatible’ polymer nanocomposites.39 These defect sites also enable MPD to further diffuse through the forming PA layer, resulting in a thicker PA layer, most clearly observed in M4 (Fig. 5e and f). Aer casting, the partially reacted TMC monomers adjacent to C60@PAF900 particles are later converted to free carboxylic acids when the TFC/TFN membranes are washed in DI water, explaining the observed trend in FITR (Table S2†). With polymerization forced to propagate around settled C60@PAF900 particles, agglomerates would be expected to be visible embedded within “valleys” or on top surfaces of the PA

Fig. 7 Influence of C60@PAF900 on the interfacial polymerization. Purple arrows show the diffusion pathways of MPD: (a) C60@PAF900 located at the interface of TMC and MPD, (b) C60@PAF900 covered or half-embedded by polyamide with low loading and (c) C60@PAF900 exposed on the polyamide surface with high loading.

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layer, matching our observations (Fig. 7c). Although disruption might be expected to be reduced by the agglomeration of particles (reducing the effective barrier area at the solvent interface within the forming polymer), it appears that higher concentrations of C60@PAF900 also accelerate the settling of particles from the organic phase, leading to more severely defective PA layers. Fig. S5† shows the water contact angles of the FO membranes as a function of time from 0 to 300 s. The initial water contact angles for all the TFN membranes were larger than that in the M0 membrane at 0 s, increasing from 85.4 (M0) to approximately 102.2 degrees (M1) with 0.005 w/v% C60@PAF900 added into the PA layer of TFN FO membranes. Higher C60@PAF900 concentrations had only a slightly greater effect on the contact angle, suggesting a similar coverage of the hydrophobic C60@PAF900 on the PA layer of each TFN membrane. Interestingly, the contact angles of all membranes decreased with time indicative of surface wetting and adsorption into the membrane. Moreover, at 30 s, the contact angles in the M3 and M4 membranes were less than that of M0, whereas aer 300 s, the water contact angles of the M2, M3 and M4 membranes were either similar or less than that of the M0 membrane (Table 1). The reduction degree in M3 was a little higher than that in M4, which was potentially due to the aggregation of C60@PAF900 when the concentration was high. In the M4 membrane, there was more C60@PAF900 exposed on the surface of the polyamide layer, without embedded by –COOH. As a result, the water molecules were not easily attached on the surface of C60@PAF900, which reduced the transportation of water molecules. Comparing the initial contact angle values with that aer 300 s, it was found that the wetting/adsorption effect for each of the TFN FO membranes was larger than that in the M0 membrane, indicating that water molecules are transported faster into/through the TFN FO membranes. The primary reason for this improvement was the inuence of C60@PAF900 on the structure of the PA layer. The inclusion of C60@PAF900 appears to cause interfacial defects in the PA layer due to the graphene repulsion, which may contribute to the increase in the overall water permeability by enabling transport through the void space present between C60@PAF900 and the polyamide matrix.39 In addition, the PA layer of the FO membranes was prepared by the interfacial polymerization process between MPD and TMC. As a result, the PA layer formed many hydrogen bonds between the amide groups and carboxylic acid moieties and subsequently with permeating water.49 Table 1 Water contact angles of the polyamide surface of the FO membranes and the reduction of water contact angles after coming into contact for 300 s

Contact angle (degree) Membrane

0s

300 s

Reduction degree (%)

M0 M1 M2 M3 M4

85 102 107 107 105

60 66 61 50 52

30 35 43 54 51


Because of the hydrogen bonds, the water transportation rate in the PA layer slowed down. As shown in the FTIR results (Fig. 4 and Table S2†), more –COOH functional groups appeared in the PA layers of the TFN FO membranes as the concentration of C60@PAF900 increased. These hydrophilic functional groups were embedded around the C60@PAF900, which could facilitate the adsorption of water molecules on the C60@PAF900 surface. C60@PAF900 lacks functional groups to form hydrogen bonds. As a result, the water molecules could transport quickly through the pores of the C60@PAF900 without the impediment of affinity between the hydrophobic pores of C60@PAF900 and water molecules. This result was in line with previous studies investigating the effect of hydrophobic porous materials.15,20,22,23 Moreover, Tunuguntla et al.23 reported that water molecules easily formed a single-chain conguration in carbon nanotube ˚ pores, leading to an enhanced water transporins with 8 A portation compared with that in carbon nanotube porins with ˚ pores. Here, the C60@PAF900 main pore size is 5 A, ˚ which 12 A while still considerably larger than the kinetic diameter of water may further enhance the transportation of water molecules by the described single-le conguration mechanism.23 Fig. 8 shows the water ux (JW) and reverse solute ux (JS) of the TFC and TFN FO membranes in the ALFS mode and the ALDS mode using 2 mol L1 NaCl and 10 mmol L1 NaCl as the draw solution and feed solution, respectively. Compared with the pristine TFC FO membrane (M0), the C60@PAF900 modied TFN FO membranes showed higher water uxes. The water ux increased from 7.4 to 12.4 LMH in the ALFS mode and from 12.6 to 21.3 LMH in the ALDS mode, with the optimal water ux of the TFN FO membranes achieved at 0.015 w/v% C60@PAF900 in the organic solution (M3). This result was in line with the result of water contact angle reduction (Table 1), which showed that the M3 membrane has the largest reduction degree during the test time. The reason for this improvement was due to the facilitation of water transportation in PA layers by adding porous C60@PAF900. The improved void space between C60@PAF900 and the polyamide matrix enhanced the water transport. Moreover, when the concentration of C60@PAF900 was lower, the C60@PAF900 was covered or half-embedded by the PA layer. As a result, C60@PAF900 particles were surrounded by the hydrophilic polyamide. When water molecules went through the membrane surface, they were attached on the polyamide surface due to the hydrogen bonds between water molecules and hydrophilic functional groups on the PA layer. Because of weaker interaction between the hydrophobic walls of pores in C60@PAF and water molecules, the water could transport through the interfacial regions within the PA layer easily compared to the bulk crosslinked PA layer in the TFC membranes. However, when the concentration of C60@PAF900 exceeded 0.015 w/v%, the water ux begun to reduce. This may be caused by the aggregation of C60@PAF900 at high concentrations, which by impeding the polymerization reaction causes the thickness of the TFN membranes to increase. Furthermore, the degree of polyamide crosslinking also decreases with increasing C60@PAF900 content, meaning that high concentrations of unreacted carboxylic acid groups are also present, which were

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Paper Table 2

Properties of the prepared FO membranes

JS/JWa Water permeability Ab Salt permeability Bc B/A Membrane (g L1) ( 1012 m s1 Pa) ( 108 m s1) (kPa) M0 M1 M2 M3 M4

0.58 0.61 0.62 0.84 1.2

3.31 4.27 6.84 8.87 7.20

5.00 6.67 11.73 18.34 18.33

15.1 15.6 17.1 20.7 25.5


a JS and JW were measured using a FO test system. 2.0 mol L1 NaCl was used as the draw solution, while 10 mmol L1 NaCl was the feed solution. b A was measured using a RO test system at 0.2 MPa and DI water was used as the feed solution. c B was measured using a RO test system at 0.2 MPa and 20 mmol L1 NaCl was used as the feed solution.

Fig. 8 Water flux and reverse solute flux of the FO membranes in (a) the ALFS mode (the active layer facing the feed solution) and (b) the ALDS mode (the active layer facing the draw solution). 2.0 mol L1 NaCl was used as the draw solution, while 10 mmol L1 NaCl was the feed solution.

previously shown to slow the transport of water in CNT membranes.21 Together, these factors offer an explanation for the deteriorated performance of the M4 membrane. As the concentration of C60@PAF900 increased, the reverse solute ux (JS) measured using both ALFS and ALDS modes increased for each of the TFN FO membranes (Fig. 8). The ratio of reverse solute ux and water ux (JS/JW) is applied to evaluate the efficiency of desalination performance of the FO membranes.50 A lower JS/JW value indicates a smaller draw solute loss in the FO ltration process and a comparatively higher efficiency FO membrane.51 Table 2 shows the JS/JW of the FO membranes in the ALFS mode. Comparing the JS/JW values of M0, M1 and M2 membranes, the JS/JW value slightly increased from 0.58 g L1 to 0.62 g L1 with the concentration of C60@PAF900 increasing from 0 to 0.01 w/v%. Table 2 also shows the water permeability (A) and solute permeability (B) of the prepared FO membranes. Compared with M0, the A was higher in each of the TFN FO membranes due to the discussed improvement of water transportation. The value of B/A showed the efficiency of the FO membranes. A higher B/A indicated a lower efficiency of the FO membranes.52 Similar to the

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tendency of JS/JW, low loading of C60@PAF900 (M0, M1 and M2) showed less inuence on the values of B/A. However, when the concentration of C60@PAF900 increased to 0.015 w/v% and 0.02 w/v%, the JS/JW and B/A increased signicantly (M3 and M4), which was potentially due to the aggregation of C60@PAF900 at these high concentrations. Fig. S6† shows the B and A values of the FO membrane fabricated in this work and four types of commercial membranes (CTA-HW, CTA-W, CTANW and BW30), which were tested by Wei and co-authors.53 The result indicated that the M1–M4 membranes showed higher A values compared with CTA-W, CTA-NW and CTA-HW membranes. Moreover, the M3 membranes showed the highest A values among all other membrane samples. The B value of the M3 membrane was less than that of the CTA-HW membrane, which has the highest A value among the commercial membranes. This result indicated that with a small addition of C60@PAF900, the water ux of TFN FO membranes can be enhanced without sacricing too much salt rejection.

Conclusions In this study, novel TFN FO membranes were fabricated by incorporating C60@PAF900 into the PA layer through the interfacial polymerization process. Compared with the pristine TFC FO membranes, the TFN FO membranes varied with differences in polyamide surface morphology, wettability and desalination performance. The crosslinked network in the polyamide surface layer of the TFN FO membranes was shown to be signicantly affected by the inclusion of C60@PAF900 and less crosslinked as the concentration of C60@PAF900 was increased. In the resulting TFN FO membranes, the PA layer became more hydrophobic due to the hydrophobic nature of C60@PAF900. However, the contact angles of the polyamide surface in the TFN FO membranes reduced more signicantly than those of the TFC FO membrane, indicative of improved water transport through the membrane and pores of C60@PAF900. More specically, the contact angle of 0.015 w/v% C60@PAF900 modied TFN FO membranes reduced from 107 to 49.5 degrees within the testing period of 300 s, while the contact angle in the TFC-FO membrane only dropped from 85.4 to 59.5 degrees. In terms of the desalination performance, the C60@PAF900 modied TFN FO membranes showed higher water


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ux. In particular, compared with the pristine TFC FO membranes, 0.015 w/v% C60@PAF900 increased the water ux up to 67% and 69% in the ALFS and ALDS modes, respectively. This study demonstrates the potential of unusual porous materials such as C60@PAF900 to inuence membrane properties toward the development of advanced FO membranes. Further work to prevent the aggregation of C60@PAF900 at high concentrations and control the formation of free acid of linear PA chains through advanced additives may prove to be a successful and novel approach to overcoming the current challenges in advanced TFN FO membrane fabrication in water purication technology.


Conflicts of interest There are no conicts to declare.

Acknowledgements This work was supported through the grants from CSIRO Manufacturing, the Bureau Frontier Sciences and Education (QYZDB-SSW-DQC044), the Bureau of International Cooperation (132C35KYSB20160018), and the Chinese Academy of Sciences and the grants from the Joint Program between CAS and CSIRO. X. Wu appreciates his anc´ ee Ms Siqian Tu for English editing, the technical help from Ms. Chen Zhao and Monash CryoEM/Ramaciotti Centre and the scholarship from the China Scholarship Council.

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