Improving Gas Separation Performance of Poly(vinylidene fluoride

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Improving gas separation performance of poly(vinylidene fluoride) based mixed matrix membranes containing metal organic frameworks by chemical modification Elahe Ahmadi Feijani, Ahmad Tavasoli, and Hossein Mahdavi Ind. Eng. Chem. Res., Just Accepted Manuscript • DOI: 10.1021/acs.iecr.5b02549 • Publication Date (Web): 16 Nov 2015 Downloaded from http://pubs.acs.org on November 18, 2015

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Industrial & Engineering Chemistry Research

Improving gas separation performance of poly(vinylidene fluoride) based mixed matrix membranes containing metal organic frameworks by chemical modification Elahe Ahmadi Feijani, Ahmad Tavasoli, Hossein Mahdavi* *

School of Chemistry, College of Science, University of Tehran, Tehran, Iran P.O. Box 14155-6455, Tel/Fax: +98-21-66495291([email protected])

Abstract In this study poly(vinylidene fluoride) (PVDF) was modified by a mixture of KOH and KMNO4 in order to HF elimination. During this reaction, some functional groups were created in PVDF structure. The metal organic frameworks, MIL53(Al) and NH2-MIL-35(Al), were embedded in modified PVDF (M-PVDF) to fabricate mixed matrix membranes (MMMs). Different characterization techniques such as FT-IR, XRD, SEM, DSC, TGA and contact angle tests have been implemented to recognize the prepared membranes. Pure and mixed (1:1) CO2 and CH4 were used to test gas separation performance of MMMs. CO2 permeance was increased for pure M-PVDF membrane compared to the pristine PVDF by 31.2%. Upon 10 wt% loading of MIL-53(Al) and NH2-MIL-35(Al) in M-PVDF, CO2 permeability increased 104.33 and 80.02% relative to the unfilled M-PVDF. The highest CO2/CH4 ideal selectivity and separation factor of 43.9 and 42.65 were reported for NH2-MIL-35(Al)/M-PVDF with 10 wt% loading, respectively. These improvements could be attributed to the breathing effect of MOFs, created polar moieties in PVDF through the modification and hydrogen bonding interactions between the MOFs and the polymer. 1- Introduction Over the last two decades, different inorganic filler particles have been dispersed in polymer matrices to fabricate the so-called mixed matrix membranes (MMMs), first reported by Zimmerman [1]. In the gas separation processes, MMMs are promising methods to overcome the main constraint of polymeric membranes which is the trade-off between permeability and selectivity [2,3]. MMMs benefit of the superior gas transport properties afforded by inorganic fillers with the ease of

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fabrication and optimal mechanical properties of polymer matrices. Appropriate choice of polymer and filler is the most important challenges in which researchers are seeking for it. Gas transport properties of MMMs strongly depend on the filler gas property and polymer–particle interface morphology. Non-selective voids, rigidification of polymer chains around particles and partial blocking of particle pores by polymer are the most common morphologies that happen at the interface region. Low compatibility between polymer and filler can result in formation of non-selective voids at the interfacial region. Gas molecules are allowed to pass through such non-selective voids with no distinct selectivity. Consequently, the gas selectivity performance of whole membrane would be reduced. Inhibition of polymer chain mobility at the interface region creates chain rigidification morphology in MMM. Chain rigidification affects the membrane separation performance in a different way compared to the non-selective voids. Partial blockage of the filler pores by penetrating of polymer chains can impress interfacial morphology. If reduced pore size caused by partial pore blocking is appropriate to occur molecular sieving mechanism, this morphology can help to improve gas separation behaviour of membrane[4]. In addition, using fillers with appropriate surface functional groups and tendency toward some gases is another important factor which should be considered in MMM fabrication. Various families of fillers such as zeolites [5,6], carbon molecular sieves [7], inorganic oxides [8], carbon nanotubes [9,10], carbon nanofiber [11], mesoporous silica [12,13], fullerenes [14], modified clay [15] and metal organic frameworks (MOFs) [16], have been utilized to incorporate into polymers. Recently MOFs have attracted attention of many researchers in MMM preparation since they constitute of transition metal complex or cluster and polyfunctional organic linker. In principle, organic ligands with broad different types of 2 ACS Paragon Plus Environment

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functionalities interact with polymer chain pretty well. Thus, formation of nonselective void at the polymer-MOF interface, which deteriorate membrane gas separation performance, can be prevented [4]. Moreover, MOFs can be prepared at moderate conditions and possess considerable characteristics such as large inner surface area, tunable pore size, uniform pore size distribution, controlled porosity, various topologies, flexibility in their chemical composition and affinity toward certain gases [4,17]. Different studies were done to investigate the effect of Al based MOFs on performance of MMMs. The effect of NH2-MIL-53 in 6FAD-ODA polyimide for gas separation process was investigated by Chen et al. The MMMs with 32 wt% of MOF showed high ideal selectivity of 77 and consequently high separation factor of 53. These results indicated the great adhesion between the polymer and NH2-MIL-53 particles due to the hydrogen bonding interactions of NH2 groups and polyimide [18]. The synthesized NH2-MIL-53 by microwave heating was embedded in Polysulfone (Udels P-3500) matrix to examine CO2/CH4 separation. Fabricating MMM with high filler up to 40 wt% loading was possible due to the strong matching of NH2MIL-53 with PSF. Optimal loading of 25 wt% of MOF increased CO2/CH4 selectivity from 25 for MOF-free PSF to 46 approximately [19]. Basu et al. in the study of mixed-matrix membranes comprising MOFs reported that MIL-53(Al) had higher selectivity for CO2 over CH4 than ZIF-8 in Matrimid matrix. The superior gas separation behavior was due to the strong CO2 interaction with the hydroxyl groups of the MIL-53(Al) and breathing mechanism of MIL53(Al) in presence of CO2 [20]. Chen et al. explored the nano-sized amino-functionalized MOF, NH2-MIL-53, with two commercial polyimides (Matrimid 5218 and Ultem 1000) for gas separation. 3 ACS Paragon Plus Environment


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Matrimid/NH2-MIL-53 MMM showed high permeability for both CO2 and CH4 gases and relatively low CO2/CH4 selectivity compared to the pure membrane. Low selectivity of Matrimid/NH2-MIL-53 was due to the non-selective void formation at the interface between polyimide and MOF particle. For Ultem/MOF, permeability of CO2 and CH4 and CO2/CH4 selectivity were higher than the neat Ultem since Ultem had better compatibility with MOF crystals [21]. In another study, the Al based functionalized MOFs were embedded in Matrimid for CO2 and CH4 separations. The separation factor about 36 was obtained for 15 wt% of NH2-MIL-101(Al). For 20 wt% of NH2-MIL-53(Al) in Matrimid, although the CO2/CH4 separation factor increased 15%, CO2 permeability was decreased slightly compared to the pure membrane [22]. The impact of NH2-MIL-53(Al) and NH2-MIL-101(Al) in sulfur-containing copolyimides matrix (6FDA:DSDA/4MPD:4,4’-SDA 1:1 (polymer P1) and 6FDA/4MPD:4,4’-SDA 1:1 (polymer P2)) was studied by Seoane et al. for gas separation process. The permeation of H2, CO2 and CH4 gases were tested through the membranes. The gas separation results for fillers with P1 were better than with P2 since the flexible SO2 groups in P1 had positive effect on the filler-polymer matrix interaction. 10 wt% of NH2-MIL-101(Al) in P1 showed high permeabilities of 114, 71 and 1.7 Barrer for H2, CO2 and CH4, respectively [23]. The performance of NH2-MIL-53(Al)/Matrimid@5218 MMM was reported by Rodenas et al. For 8 wt% loading of MOF, CO2 permeability was lower than the pure polymer. With further MOF loading up to 25 wt%, a 50% increment in CO2 permeability was observed with respect to the unfilled membrane. CO2/CH4 separation factor increased in 8 wt% loading while retained unchanged in 25 wt% loading [24].


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Different types of fillers (ZIF-8, NH2-MIL-53, MCM-14, ETS-10 and silica) were used to fabricate MMMs with Polysulfone (PSF) matrix by spin coating method. Over 60% enhancement in H2/CH4 and O2/N2 separations was observed for MMM contained NH2-MIL-53 compared to the pure polymer. Good affinity of NH2-MIL53 toward PSF chains was responsible for better gas separation behavior of NH2MIL-53 among the other fillers [25]. Outstanding properties such as excellent mechanical strength, thermal stability, chemical resistance and high capability of film forming have made PVDF interesting for membrane industrial applications such as membrane distillation [26] and membrane contactor [27]. Generally PVDF has been recognized as a gas barrier material and its application in gas separation membranes is in less importance. However, incorporating appropriate fillers into the PVDF matrix accompany with PVDF modifying can provide opportunity to use this polymer for gas separation applications. Although PVDF is a semicrystalline polymer, it has flexible chains which is expected to form an intimate interaction with fillers like MOFs. MIL-53(Al) materials were chosen to fabricate MMMs with PVDF. MIL53(Al) is build up from AlO4(OH)2 octahedra infinite chains cross linked by means of benzene dicarboxylate organic ligands. The three-dimensional microporous framework of MIL-53(Al) contained unidirectional diamond shaped channels with internal diameter of 0.85 nm [28]. NH2-MIL-53(Al) has an analogue structure based on MIL-53(Al) topology [29]. Because of the flexible bond between Al3+ metal centers and organic ligands, both MIL-53(Al) and NH2-MIL-53(Al) exhibit special breathing behavior. However, their breathing behavior is different. Molecules which possess dipole or quadrupole moments such as CO2 and H2O can induce breathing effect. In the absence of guest molecules, large pore configuration is energetically dominated by 3 kJ/mol-1 the while strong interaction of guest 5 ACS Paragon Plus Environment


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molecules in low pressure with hydroxyl group of MIL-53(Al) contracts pore size into the narrow pore configuration. The opposite breathing behavior is reported for the amino functionalized MIL-53(Al). Unlike MIL-53(Al), for NH2-MIL-53(Al) the narrow pore configuration is more stable than large pore by 14 kJ/mol-1 at the absence and low pressure of CO2. It expands to its large pore configuration at high pressure of CO2 [30]. In addition, high surface area, large pores and chemical resistance of MIL-53 series are also responsible for their excellent tendency for the selective adsorption of CO2. In the case of NH2-MIL-53(Al), presence of amino functional groups increases both CO2 adsorption and interaction with existing functional groups in polymer matrix which result in high CO2/CH4 separation [18]. In the present study, poly(vinylidene fluoride) (PVDF) was chemically modified by using KMnO4 and KOH as oxidant and strong base. The chemical modification led to create some functional groups in polymer chains. Then the modified PVDF (M-PVDF) was used to fabricate MMMs with MIL-53(Al) and NH2-MIL-53(Al). The impact of chemical modification and MOF loadings on CO2/CH4 separation properties of PVDF was evaluated by measuring the gas permeability and selectivity.

2- Experimental 2-1- Materials N,N-dimethylacetamide (DMAc), N,N-dimethylformamide (DMF), acetone, Al(NO3)3.9H2O, 1,4-benzenedicarboxylic acid and 2-amino-1,4benzenedicarboxylic acid (NH2-H2BDC), KOH and KMnO4 were received from Sigma Aldrich. 2-2- Modification of PVDF PVDF was modified according to a procedure which was published previously [31]. PVDF powder was poured to the aqueous solution of KOH (0.5 mol/l) and KMnO4 (1.0%). The reaction was kept at 60 ºC for 10 minutes. The color of PVDF 6 ACS Paragon Plus Environment

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altered from white to brown. The brown precipitate was filtered and moved to the acidic solution of 1.2% NaHSO3 until it regained white color. Finally modified PVDF was washed 2 or 3 times with distilled water and dried at room temperature under vacuum for 48h. 2-3- MOF preparation The MIL-53(Al) and NH2-MIL-53(Al) MOFs were synthesized according to the method described in previous study [32]. In synthesis of MIL-53(Al), measured amounts of Al(NO3)3.9H2O (0.375 g) and 1,4-benzenedicarboxylic acid (0.268 g) were added to DMF (14 ml) under stirring. The solution was moved to a Teflon lined steel autoclave in an oven at 130 ºC for 3 days. The obtained white crystals were washed with DMF to remove unreacted reagents. Soxhlet extractor was utilized to replace DMF with a solvent with a low boiling point like acetone. The result crystals were dried at 100 ºC under vacuum for 24h. The same procedure was used to synthesize NH2-MIL-53(Al) where 1,4-benzenedicarboxylic acid was replaced by 2-aminoterephthalic acid. 2-4- Preparation of M-PVDF dense membrane The dense film casting method was applied to fabricate membranes. The polymer powder was kept under vacuum at 100 ºC for 24h to remove any adsorbed moisture. Then the solution of 10% polymer in DMAc was prepared at room temperature during 48h. Thereafter, the solution was degassed to draw out any bubble which created during stirring. After complete degassing, the solution was cast onto glass substrate and dried at 80 ºC in an oven for 12h. At last, the membrane was kept in a vacuum oven for 24h at 120 ºC in order to remove any remnant solvent. 2-5- Preparation of mixed matrix membranes All of the mixed matrix membranes were prepared as follows. The measured amount of MOF was added to DMAc solvent and the mixture was stirred for 24h. In addition, bath sonication was used for 20 min to make sure that MOF was homogenously dispersed in solvent. The polymer was added to the MOF solution by using priming technique. Priming technique can enhance adhesion between the polymer and MOF so improves gas separation properties of membranes. In this regard 10% of polymer was added to the mixture and the solution was mixed for 5h. Thereafter the remaining amount of polymer was added and stirring was 7 ACS Paragon Plus Environment


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continued to obtain a uniform mixture. After degassing the solution, it was cast onto glass substrate and dried with the same treatment which was used for the dense membrane. Mixed matrix membranes comprising 5%, 10% of MIL-53(Al) and NH2-MIL-53(Al) in M-PVDF were fabricated to compare their gas separation efficiency. The required amounts of MOFs were calculated based on the following equation: wt.MOF   MOFloading =   × 100  wt.MOF + wt. polymer 

2-6- Gas permeation measurement Permeability determination of pure gases was implemented by using the variable pressure/constant volume method for all of the membranes. This method is appropriate for membranes with low permeability. The whole system was evacuated for 4 h before every measurement. The pressure of the penetrated gas in the upstream chamber was fixed at desired amount. A high precise pressure transducer (0-100 mbar, BD sensors, Germany) was used to monitor pressure change in downstream chamber. Permeability and ideal selectivity were calculated using the following equations [33]: P( Barrer ) =

273.15 ×1010Vt dP ( ) 760 × 76 ATP dt

S electivity =

Pi Pj

where t is membrane thickness (cm), V is the volume of the downstream chamber (cm3), A is the membrane surface area (cm2), T is the experimental temperature (K), P is feed pressure, and dP/dt is the pressure gradient (mmHg/s) in downstream chamber. In this study i and j are CO2 and CH4 gases and Pi and Pj are defined as CO2 and CH4 permeabilities, respectively. The following equation was used to calculate the separation factor (αi/j) of membranes in a binary gas mixture of CO2 and CH4 with ratio of 1:1. αi/ j =

yi / y j xi / x j


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where x and y are mole fractions of gases in feed and permeate, respectively. The composition of feed and permeate in mixed gas test were determined by GC analysis. After reaching the steady state, at least four measurements were done at different time and the average values were reported. 2-7- MOF and membrane characterization Field emission scanning electron microscopy (FESEM) ( HITACHI, Model S4160) was used to study morphology of the MOFs and particle–polymer interface. The membranes were broken in liquid nitrogen and covered a thin layer of Au. A differential scanning calorimeter (Netzsch DSC 200F3) has been used to determine the melting temperature of selected membranes. The membranes were kept at 40 ºC and then temperature increased to 200 ºC with a heating rate of 5 ºC/min under nitrogen atmosphere. The melting temperatures (Tm) and fusion heat ( ∆H ) of the selected membranes were reported. Powder X-ray diffraction patterns of the prepared membranes were achieved by using an X-ray diffraction instrument (Siemens D5000) with a Cu target ( λ=0.1540 nm) at RT over a 2θ range from 0 to 50º. Thermal resistance of the prepared membranes was investigated by PL termogravimetric analyzer (TGA 1000, UK). The TGA curves were obtained over the temperature range of 30 to 800 ºC with a heating rate of 10 ºC/min under nitrogen stream. Physisorption of nitrogen for MOFs was done by Belsorp mini II, BelJapan instrument at 77 k. Surface area of the samples was calculated by the Brunauer-Emmet-Teller (BET) method. Mechanical properties of the membranes were evaluated by a Z030 Zwic/Roell testing machine at room temperature with an operating rate of 1 mm/ºC. FT-IR spectra of MOFs were recorded with Bruker FTIR specterometer. 3- Results and discussion 3-1- MOF characterization The MOF were characterized by nitrogen physisorption–desorption, XRD, TGA and FT-IR which have been mentioned extensively in previous study [32]. The cryogenic nitrogen physisorption–desorption results of MOFs are shown in Table 1 which were in agreement with the values reported in literature.



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Table 1: Characteristics of the synthesized MOFs

MOF

structure Pore channel BET surface area (m2/g) MIL-53 3D 1D 1130.19 1118 1140 1365 1408 1170 NH2-MIL-53 3D 1D 735 795 780 782

Pore volume (cm3/g) 0.56 0.433 0.462 0.424 0.240 0.258 0.246

Ref. [34] [35] [36] [18] [33] This work [18] [33] [37] This work

The morphology of MIL-53(Al) was not changed by introducing the functional group however the BET surface area and the pore volume were influenced. The average particle sizes of MOFs were estimated around 100 nm. Thermal resistance of MIL-53(Al) and NH2-MIL-53(Al) were obtained up to 500 and 420 ºC, respectively. The lower stability of the functionalized one was clearly due to the existence of amino functional groups in MOF structure. 3-2- Membrane characterizations The interaction between M-PVDF and MOFs were considered by FT-IR technique. FT-IR spectra of the prepared membranes and MOFs are depicted in Fig. 1. From the comparison of PVDF FT-IR before and after modification, the appeared absorption peaks at 1703 and 3200-3600 cm-1 indicated that modifying of PVDF was accomplished successfully. The former was related to the carbonyl group and the latter showed hydroxyl group existence in M-PVDF. In the first step of modification, HF was eliminated and during the nucleophilic addition reaction, the OH groups were created in PVDF chains. Existence of C=O groups might be due to the oxidation of modified polymer in the air [31]. It should be noted that the peaks at 2900 cm-1 was attributed to the CH2 group. In the case of pure MOFs, characteristic peaks at 1511 and 1597 cm−1 were assigned to asymmetric stretching vibration of COO and the absorption peak at 1416 cm−1 was corresponded to symmetric stretching vibration of COO. The absorption peak at 3380 and 3490 10 ACS Paragon Plus Environment

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cm−1 in functionalized MOF were attributed to NH2 asymmetric and symmetric stretching vibration. Formation of hydrogen bonds between the M-PVDF and MOFs could be the reason for good interaction of both phases. The observed shift in characteristic peaks related to the MOFs or polymer confirmed hydrogen bonds formation which has been reported by other researchers [6,21,33]. The peak related to the carbonyl of M-PVDF was shifted to 1762 and 1790 cm−1 for 10% loaded MPVDF with MIL-53(Al) and NH2-MIL-53(Al), respectively. The three peaks attributed to the C=O of MOFs were transferred to other wavenumbers for the prepared MMMs. NH2 absorption bands were located in lower frequency for the MMM compared to the pure NH2-MIL-53(Al). NH2 absorption bands appeared at 3373 and 3488 cm−1 for the MMM containing 10% of the functionalized MOF.

Fig. 1: FT-IR spectra of a) MIL-53(Al) b) NH2-MIL-53(Al) C) PVDF d) M-PVDF e) MPVDF/MIL-53(Al)-10% f) M-PVDF/NH2-MIL-53(Al)-10%

Membrane’s surface property was determined by contact angle test. Table 2 shows the effect of modification and MOF loading on variation of water contact angle of 11 ACS Paragon Plus Environment


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PVDF. Pure PVDF membrane showed a contact angle of 89º. After modification of PVDF, the contact angle decreased to 56º indicating the formation of hydrophilic groups in PVDF as was proved by FT-IR analysis. The hydrophilicity of the membranes increased at the presence of MOF particles due to the hydrophilic nature of the fillers with surface OH groups. The contact angle of MMMs containing 10 wt% of MIL-53(Al) and NH2-MIL-53(Al) was decreased by 48.3 and 49.4% as compared with pure PVDF membrane. Table 2: Contact angle of prepared membranes

MOF loading 0 5 10

MIL-53(Al) 66.7±0.4 56.5±0.6 46.8±0.3

NH2-MIL-53(Al) 66.7±0.4 54.4±0.2 45.3±0.3

SEM images of pure M-PVDF membrane and 10 wt% M-PVDF/MOFs MMMs are given in Fig. 2. Dispersion of the MOF particles was uniform throughout the polymer matrix. The X-ray mapping of membrane cross sections accordingly showed good particle dispersion as well (Fig. 3). No “sieve-in-a-cage” morphology was detected in SEM images and matching of MOFs and polymer matrix were excellent. In addition, increase in plastic deformation of polymer matrix in MMMs compared to the pure membrane could confirm the complimentary contact between particles and M-PVDF. Plastic deformation of polymer matrix was due to the interfacial stress during cryogenically fracturing membranes in liquid nitrogen. The interfacial stress was created because of the strong interaction between phases [38,39].


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Fig. 2: SEM images of prepared membranes, a) pure M-PVDF, b) M-PVDF/ MIL-53(Al), c) M-PVDF/NH2-MIL-53(Al) with two different magnifications



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Fig. 3: X-ray mapping of the cross section of a) PVDF/MIL-53(Al) and b) PVDF/NH2-MIL53(Al)

To determine the crystal structure of membranes, XRD analysis was performed. Four common crystal phase (α, β, ϒ, δ) have been reported for PVDF. Crystallinity of PVDF is one of the most important factors which influence its mechanical properties. Meanwhile the non-polar α phase has been known for the most common crystal phase of PVDF. With exposure of α phase to different conditions like mechanical stretching, electric field, high-temperature, etc. other crystal phases can be obtained. So crystal content of PVDF membranes based on experienced conditions is different. There have been studies to investigate the effects of variables on crystalline structure of PVDF membranes [40-43]. The XRD patterns of pure membranes and M-PVDF/MOFs MMMs are shown in Fig. 4. The observed characteristic peaks at θ= 17.06, 18.3 and 20.17º for M-PVDF was similar to the original PVDF membrane. These peaks were assigned to (1 0 0), (0 2 0) and (1 1 0) lattice planes of α phase, respectively. The β phase was absent in XRD patterns of pure membranes since the related peak did not emerge in spectrum. This demonstrated that the dominant crystal phase in PVDF or the modified one was α phase. No considerable change in phase crystal of polymer matrix was observed by embedment of MOFs in M-PVDF. The pattern of MMMs exhibited the peaks related to the crystalline structure of MIL-53(Al) and NH2MIL-53(Al). Lack of matching the XRD pattern of NH2-MIL-53 based MMM with NH2-MIL-53 particle was due to the pore configuration of MOF particle. The flexible framework of the MOF was changed from narrow pore form after 14 ACS Paragon Plus Environment

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synthesis to its large pore configuration in MMM. This configuration variation could be attributed to partial penetration of polymer chains in the pore structure of the NH2-MIL-53(Al). The same phenomenon was reported at previous work and other literature [24]. So the ordered crystalline structure of MOFs was maintained in MMMs during membrane preparation.

Fig. 4: XRD patterns of a) MIL-53(Al), b) NH2-MIL-53(Al), c) pure PVDF, d) pure MPVDF membrane and MMMs containing 10 wt% of e) MIL-53(Al) and f) NH2-MIL-53(Al)

Fig. 5 presents TGA curves of the prepared membrane. Thermal stability of pure M-PVDF membrane decreased slightly compared to the pure PVDF. This could be due to the introduction of hydroxyl group in PVDF through the modification process. Although PVDF possesses excellent chemical resistance, alkaline exposure of PVDF diminished its stability considerably. Despite the good interaction of MOFs and M-PVDF, thermal stability of their MMMs decreased substantially. Degradation temperature (Td) of pure M-PVDF and 10% loaded of M-PVDF with MIL-53(Al) and NH2-MIL-53(Al) were reported as 470, 386 and 386 ºC, respectively. As was mentioned in previous work, lower thermal stability 15 ACS Paragon Plus Environment


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of MOFs relative to the polymer matrix, thermal degradation products of MOFs and M-PVDF and interaction of HF with these degradation products could be responsible for such considerable decrease in the decomposition temperature of MMMs. It is worthy to note that Td of 10% loaded of M-PVDF with MIL-53(Al) and NH2-MIL-53(Al) were slightly higher than that of MMMs with PVDF. Affinity of MOFs toward M-PVDF was probably stronger than with PVDF due to the presence of carbonyl and hydroxyl groups in M-PVDF structure.

Fig. 5: TGA curves of the prepared membranes

The DSC curves of pure PVDF, M-PVDF and MMMs containing 10 wt% of MIL53(Al) and NH2-MIL-53(Al) are illustrated in Fig. 6. Melting temperature (Tm), fusion heat ( ∆H ) and the crystallinity ( χ) of the selected membranes are given in


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Table 3. The following equation was applied to calculate the crystallinity of membranes [44,45]. χ=

∆H W × ∆H

°

× 100%

where ∆H º is the fusion heat and W is the polymer content in the MMMs. A value of 90.4 J/g was used for PVDF with 100% crystallinity of α phase.

Fig. 6: DSC curves of the prepared membranes a) PVDF b) M-PVDF c) M-PVDF/MIL53(Al) d) M-PVDF/NH2- MIL-53(Al)

Table 3: Melting temperature (Tm), fusion heat ( ∆H ) and crystalinity ( χ ) of the prepared membranes

Membrane PVDF M-PVDF PVDF-MIL-53(Al)-10% PVDF-NH2-MIL-53(Al)-10% M-PVDF-MIL-53(Al)-10% M-PVDF-NH2-MIL-53(Al)-10%

Tm (ºC) 162.6 163.2 163.3 162.7 163.1 163.0

∆H J/g 27.60 29.85 25.5 36.6 39.88 41.45

χ

34 36.68 31 45 49.01 50.94

Pure M-PVDF showed higher fusion heat and crystallinity. For pure M-PVDF Tm, ∆H and χ were determined as 163.2, 29.85 and 36.68% respectively. Melting temperature of M-PVDF and its MMMs with MIL-53(Al) and NH2-MIL-53(Al) 17 ACS Paragon Plus Environment


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were around 163 ºC similar to that of PVDF and PVDF/MOFs. Incorporation of MOFs into the M-PVDF increased crystallinity of the matrix up to 49.01 and 50.9%. The MOFs might play nucleation agent role during the crystallization of polymer due to the hydrogen bond formation between two phases. In addition, thermogram of M-PVDF/MOF MMMs showed two peaks. While the first small peak was related to the melting of trans-crystalline layer between MOF particles and polymer matrix, the melting of the semicrystalline polymer matrix created the second peak. Similar results were reported by Xu et al. [46] for PVDF/carbon nanotube (CNT) composites. Xu et al. reported DCS curves with two peaks for PVDF with different CNT contents. Moreover, crystallinity of M-PVDF/MOF MMMs was higher than the MMMs of PVDF and the same MOFs due to the strong M-PVDF/MOF interaction. This increscent was more pronounced for nonfunctionalized MOF than the functionalized one. This could confirm that the modification of PVDF influenced the interfacial region considerably. Tensile tests were used to investigate the impact of MOFs on mechanical properties of the membranes. The Young’s modulus, tensile strength and elongation at break of pure membrane and MMMs with 5 and 10 wt% of MOF are listed in Table 4. Young’s modulus of MMMs was improved compared to the pure PVDF and M-PVDF. Increase in crystallinity of membranes by incorporating MOFs should be considered in Young’s modulus improvement. So, higher Young’s modulus of M-PVDF containing both MOFs relative to PVDF/MOFs confirmed the positive effects of modification process on properties of membranes. Both tensile strength and elongation at break of MMMs decreased compared to the pure membranes. This downward trend indicated that flexibility of membranes have been reduced due to the presence of MOFs in polymer matrix. Table 4: Mechanical properties of the prepared membranes Membrane PVDF M-PVDF M-PVDF-MIL-53(Al)- 5% M-PVDF-MIL-53(Al)-10% M-PVDF-NH2-MIL-53(Al)-5% M-PVDF-NH2-MIL-53(Al)-10%

Tensile strength (MPa) 46.12 44.92 35.54 28.23 35.96 29.33

Elongation at break (%) 9.27 8.74 7.81 6.95 8.34 7.29


Young’s modulus (GPa) 0.711 0.691 0.773 0.886 0.801 0.895

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3-3- Gas transport properties of membranes 3-3-1- Gas permeability, ideal selectivity and separation factor of membranes Pure CO2 and CH4 measurements of prepared membranes, i.e. permeability and ideal selectivity were performed at 298 K and a feed pressure of 5 bar. At least two membranes were fabricated for each composition and three permeation measurements were performed for each membrane. The average values were reported as the final results. For all membranes, CO2 permeation was considerably enhanced (Table 5). For the pure modified PVDF, CO2 permeability increased while CH4 permeability was the same compared to the pure PVDF membrane. So, about 31.3 and 26.3% increscent in CO2/CH4 selectivity and separation factor were reported for the modified PVDF, respectively. This improvement could be due to the strong interaction between polar moieties in M-PVDF and CO2 gas with quadrapole moment. As it was shown in FT-IR (Fig. 1), carbonyl and hydroxyl groups were created during the modification process of PVDF. In addition of M-PVDF backbone impact on CO2 transport, carbonyl and hydroxyl groups were in good contact with the MOF particles. Table 5: Gas permeability, ideal selectivity and (50/50) CO2/CH4 mixed gas separation factor (α) of membranes measured at 298 K and a feed pressure of 5 bar Permeability (Barrer) CO2 CH4

Sample

Pure PVDF Pure M-PVDF M-PVDF-MIL-53(Al)- 5% M-PVDF-MIL-53(Al) -10% M-PVDF-NH2-MIL-53(Al)-5% M-PVDF-NH2-MIL-53(Al)-10%

0.915±0.121 1.201±0.208 1.752±0.169 2.454±0.306 1.693±0.224 2.240±0.391

0.043±0.011 0.043±0.009 0.049±0.008 0.062±0.008 0.045±0.003 0.051±0.010

Ideal selectivity CO2/CH4 21.27±1.66 27.93±1.34 35.75±1.76 39.58±1.04 37.62±1.55 43.92±1.59

Separation factor(α) CO2/CH4 (50/50) 20.50±0.93 25.89±1.05 35.01±1.31 36.78±0.97 36.88±1.01 42.65±1.51

Incorporation 5 and 10 wt% of MIL-53(Al) in M-PVDF increased permeability of pure gases, CO2/CH4 selectivity and separation factor. With 5 wt% of MIL-53(Al), 45.8% increment for CO2 permeability was observed relative to the pure M-PVDF. This increment enhanced up to 104.3% for 10 wt% loading of MIL-53(Al). 19 ACS Paragon Plus Environment


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Significant increase in CO2 permeation could be attributed to the polymer matrix and MOF structure. CO2 gas has high affinity toward the existence hydroxyl groups of M-PVDF and MOF structure. Because of special breathing effect of MIL-53(Al), pore configuration could be changed in presence of guest molecules like CO2. In fact adsorption of CO2 contracts the pores of MOF due to the interaction between CO2 and the hydroxyl group of MIL-53(Al). So, the configuration of MOF pores varies from large form to narrow form during CO2 adsorption. Ideal CO2/CH4 selectivity increased 28 and 41.7% for 5 and 10 wt% of MIL53(Al) compared to the pure M-PVDF membrane. The same increment trend was for separation factor, too. This increase in CO2/CH4 selectivity and separation factor was even more than the PVDF/MIL-53(Al) MMMs in previous work. The obtained results indicated that non-selective voids between MIL-53(Al) and PVDF disappeared when PVDF was modified. Lack of sufficient affinity of polymer toward the MOF leads to the non-selective void formation at the interfacial phase. Such voids deteriorate performance of the MMMs. Chemical modification of PVDF increase its compatibility with the MOF. Good compatibility was the result of hydrogen bonding formation between two phases, polymer and MOF. In addition, unoccupied metal ions in MOFs could be coordinated to the carbonyl groups in M-PVDF. This good interaction was confirmed by FT-IR analysis. Moreover, lack of interaction between CH4 and membrane accompany with reduction in pore size of MIL-53(Al) due to the breathing effect limited CH4 transport through membrane. Embedment of NH2-MIL-53(Al) in M-PVDF improved the gas permeability, ideal selectivity and separation factor. In addition of polymer matrix polar groups, the NH2-MIL-53(Al) particles played an important role in superior CO2 permeation. Introduction 5 wt% of NH2-MIL-53(Al) in M-PVDF showed 41 and 53% CO2 permeation improvement with respect to the pure M-PVDF and NH2-MIL53(Al)/PVDF MMM with the same loading. With further loading to 10 wt%, the CO2 permeation increment increased up to 86.5 and 59.3% respectively. Reversible reaction of NH2 groups in MOF structure with CO2 result in carbamate formation which favored CO2 adsorption and permeation. CO2/CH4 ideal selectivity and separation factor of NH2-MIL-53(Al)/M-PVDF MMMs substantially increased not only with respect to the pure M-PVDF, but also the NH2-MIL-53(Al)/PVDF MMMs. Besides of the good compatibility between 20 ACS Paragon Plus Environment

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the particles and M-PVDF, some polymer chain rigidification might occur around particles. Chain rigidification was the result of movement inhibition of M-PVDF in polymer-MOF interfacial area. This occurrence distinguished the penetrated gases with different molecule size. Thus, for NH2-MIL-53(Al)/M-PVDF, the CO2 permeation enhanced while CH4 decreased compared to the NH2-MIL53(Al)/PVDF and consequently CO2/CH4 ideal selectivity and separation factor significantly increased. Comparing NH2-MIL-53/M-PVDF with MIL-53/M-PVDF, the gas permeability dropped and selectivity increased. Lower gas permeability of MMMs containing NH2-MIL-53(Al) compared to those of with the non-functionalized MOF was attributed to the larger pore size of MIL-53(Al) than NH2-MIL-53(Al). Since pore size of the MIL-53(Al) (7–8˚ A) is much larger than the kinetic diameters of CO2 (3.3 ˚A) and CH4 (3.8 ˚A), both CO2 and CH4 gases can pass through the MOF pores [35]. On the contrary, smaller pore size of NH2-MIL-53(Al) led to the less gas transport. Higher CO2/CH4 ideal selectivity and separation factor of NH2-MIL53(Al) could be relevant to the better compatibility with M-PVDF and strong affinity of NH2 toward CO2. The obtained results and related results from the literature are depicted in Fig. 7. As it can be seen the performance of M-PVDF based mixed matrix membranes improved significantly compared to those with non-modified PVDF. Furthermore, with modification of PVDF, the results can be comparable with other results from literature [2,39,47-50].



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Fig. 7: Robeson curve containing performance of the MMMs in this work and in literature

3-3-2- Effect of feed pressure

The effect of feed pressure on gas performance was explored for pure M-PVDF, M-PVDF/MIL-53(Al)-10% and M-PVDF/NH2-MIL-53(Al)-10%. The difference between performance of pure membrane and MMMs at elevated pressure is depicted in Fig. 8. With increasing pressure, CO2 permeability and separation factor of pure membrane decreased as reported in the literature [18,51]. For MMMs, CO2 permeability dropped and separation factor enhanced with increasing pressure. The same trends were reported by using NH2-MIL-53(Al) and MIL53(Al) MOFs in polymer matrix in the literature [18,19,52]. The observed gas transport of M-PVDF/NH2-MIL-53(Al)-10% at higher pressure could be explained 22 ACS Paragon Plus Environment

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by MOF gas transport property. Although, increase in feed pressure increases CO2 adsorption capacity of MOF, but it does not involve increase in CO2 diffusivity through the MOF. Thus increase in separation factor suggests that strong interaction of CO2 with MOF at higher pressure excludes CH4 molecules from the MOF. Reduction in CH4 accompany with the selective CO2 adsorption by MOF lead to increase in separation factor of MMMs.

Fig. 8: Separation performance of membranes for CO2/CH4 (50/50) gas mixture as a function of feed pressure at 298 K

3-3-3- Comparing the experimental dada of MMMs with the prediction models

Most of the theoretical models for prediction of gas transport of MMMs are adapted from thermal/electrical conductivity models due to the resemblance between gas permeability through MMMs and thermal/electrical conductivity in composites. Different permeation models have been used to predict gas separation behaviour of MMMs with both interfacial ideal and non-ideal morphologies. Maxwell [53], Bruggeman [54], Lewis-Nielsen [55] and Pal [56] are the common models to describe the ability of a MMM ( Peff ) with an ideal morphology to separate gases. The effective steady-state permeability of gaseous species through



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MMM based on the Maxwell model can be calculated by using the following equation: 1 + 2φ d (α − 1) / (α + 2)

Peff = Pc

1 − φ d (α − 1) / (α + 2)

where is the Pd / Pc

φd

( Pc and

Pd

volume fraction of dispersed phase and α is the permeability ratio, are continuous and dispersed phase permeabilities respectively).

While Maxwell model is applicable for low loadings ( φ < 0.2), the Bruggeman d

model predicts permeability of gaseous species for a larger range of

φd .

However

both models suffer from the same limitations in which particle size distribution, particle shape and particle aggregation are ignored in these two models.  α −1 

( Pr )   = (1 − φd ) α −P 1/3



r

−1



where, Pr = Peff / Pc . In the Lewis–Nielsen model, the effect of particle morphology on permeability behavior of MMMs is considered by using φ parameter (the maximum packing m

volume fraction of filler particles). The Lewis-Nielsen is reduced to the Maxwell model when φ → 1 . m

Peff = Pc

1 + 2φd (α − 1) / (α + 2) 1 − ψφd (α − 1) / (α + 2)

where:  1 − φm 2  φm

ψ = 1+

  φd 

The Pal model reduces to the Bruggeman model at φ

m

 α −1 ( Pr )   α − Pr 1/ 3

  φd   = 1 −    φm 

→ 1.

− φm

Comparing experimental data with the calculated result obtained from prediction models is a useful tool for considering the interfacial phase morphology of polymer and ﬁller particles in a MMM as an ideal or non-ideal morphology. In this regard, CO2 permeation of the prepared MMMs was calculated with the four above mentioned models, Maxwell, Bruggeman, Nielsen–Lewis and Pal. The obtained results are shown in Fig. 9 and listed in Table 6. The average absolute relative errors (%AARE) were calculated to estimate the gas permeability values of pure


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MOFs. The minimum difference between experimental and prediction results give the best values of CO2 permeabilities of the pure MOFs. %AARE =

Pi cal − Pi exp  ∑ P exp  NDP i =1 i 100

NDP

where, NDP is the number of data points. Less than 1% of AARE values for both MOFs (Table 7) shows that there is a good agreement between the experimental data and results of predictive models. SEM and FT-IR results confirmed the good matching between both MOFs and polymer. Polymer matrix polar groups form hydrogen bonding with the OH or NH2 groups of the MOFs. This bond formation prevents the creation of non-selective voids in the interfacial phase. In addition of interfacial interaction, fair distribution of MOF particles in M-PVDF matrix can be concluded from the experimental and theoretical result matching, too. Table 6: Gas permeabilities and ideal selectivities of MMMs containing MOFs predicted by the Maxwell, Bruggeman, Lewis-Nielsen and Pal models MOF

MIL-53

Gas permeability and selectivity PCO2

PCO2/PCH4 NH2MIL-53

PCO2

PCO2/PCH4

MOF wt %

Experiment

Maxwell

Bruggeman

LewisNielson

Pal

18 32 1 18 32 18 32 1 18 32

1.752 2.454 35.75 39.58 1.693 2.240 37.62 43.92

1.781 2.452 11.3 36.35 39.55 1.720 2.241 9.6 38.22 43.94

1.786 2.453 10.2 36.45 39.56 1.703 2.239 7.5 37.84 43.90

1.765 2.450 9.8 36.02 39.52 1.684 2.242 7.0 39.22 43.96

1.754 2.456 7.9 35.79 39.61 1.674 2.238 6.1 37.2 43.88



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Fig. 9: Comparing the experimental data with results from the Maxwell, Bruggeman, Lewis-Nielsen and Pal models for MMMs containing MOFs

Table 7: %AARE of prediction models Prediction models %AARE Maxwell

CO2 gas MIL-53(Al) NH2- MIL53(Al) 0.87 0.82

Bruggeman

0.99

0.32

Lewis-Nielsen

0.45

0.31

Pal

0.10

0.61

4- Conclusion poly(vinylidene fluoride) was chemically modified by using KOH and KMNO4. According to the FT-IR analysis and contact angle measurements, the carbonyl and hydroxyl groups were created in PVDF structure during the modification. Increase in hydrophilicity of PVDF led to decrease in contact angle. Thus, lower contact 26 ACS Paragon Plus Environment

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angle of M-PVDF relative to the original PVDF indicated that modification of PVDF was successful. In addition, incorporation of MOFs in M-PVDF caused the decrease in contact angle of MMMs due to the existence surface hydrophilic groups in MOFs. Crystallinity of M-PVDF was enhanced with incorporation of MIL-53(Al) and NH2-MIL-53(Al) while melting temperature did not vary significantly. Thermal stability of M-PVDF/MOF MMMs decreased significantly relative to the M-PVDF. However this stability was slightly more than the PVDF/MOF MMMs. The same α crystal phase of PVDF was shown in XRD pattern for M-PVDF and incorporation of MOFs in M-PVDF did not change the α crystal phase in polymer matrix. Thermograms of MMMs showed two peaks: the small one was related to melting of the transcrystalline layer between particles and polymer and the sharp peak was due to the melting of the semicrystalline polymer. The SEM images revealed no detectable voids between MOF and matrix and the particles were dispersed in polymer uniformly. Increase in Young’s modulus of MMMs compared to pure membrane showed the positive effects of modification process on mechanical properties of membranes. The shift in some characteristic peaks in FT-IR spectra of MMMs was observed compared to the pure MOFs and M-PVDF. This confirmed that the MOFs were in a good interaction with M-PVDF matrix. Clearly, both CO2 permeability and CO2/CH4 selectivity were increased by modification on PVDF. This enhancement improved with incorporation of the MOFs with excellent affinity toward CO2. Increase in CO2 permeation was due to the interaction of CO2 with MOFs and polymer matrix. High observed CO2/CH4 selectivity confirmed excellent matching of MOFs and M-PVDF. Thus, chemical modification process not only improves gas separation performance of PVDF significantly but also did not deteriorate its chemical and mechanical properties. References 1. Zimmerman, C.M. Advanced gas separation membrane materials: Hyper rigid polymers and molecular sieve-polymer mixed matrices. Ph.D. Dissertation, The University of Texas at Austin, Austin, TX, 1998. 2. Robeson, L.M. Correlation of separation factor versus permeability for polymeric membranes. J. Membr. Sci. 1991, 62, 165–185.



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Improving Gas Separation Performance of Poly(vinylidene fluoride

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