thermal properties and structure of electrospun blends of

Thermal Properties and Structure of Electrospun Blends of PVDF

with a Fluorinated Copolymer

Nelaka Govinna ,1 Ilin Sadeghi ,2 Ayse Asatekin ,2 Peggy Cebe 1

1Department of Physics and Astronomy, Center for Nanoscopic Physics, Tufts University, 574 Boston Avenue, Medford,

Massachusetts 021552Department of Chemical and Biological Engineering, Science and Technology Center, 4 Colby Street, Medford,

Massachusetts 02155

Correspondence to: P. Cebe (E-mail: [email protected])

Received 12 November 2018; accepted 8 January 2019; published online 23 January 2019

DOI: 10.1002/polb.24786

ABSTRACT: We report the structure and thermal properties of

blends comprising poly(vinylidene fluoride) (PVDF) and a random

fluorinated copolymer (FCP) of poly(methyl methacrylate)-

random-1H,1H,2H,2H-perfluorodecyl methacrylate, promisingmem-

brane materials for oil–water separation. The roles of processing

method and copolymer content on structure and properties were

studied for fibrous membranes and films with varying composi-

tions. Bead-free, nonwoven fibrous membranes were obtained

by electrospinning. Fiber diameters ranged from 0.4 to 1.9 μm,

and thinner fibers were obtained for PVDF content >80%. As

copolymer content increased, degree of crystallinity and onset of

degradation for each blend decreased. Processing conditions

have a greater impact on the crystallographic phase of PVDF

than copolymer content. Fibers have polar beta phase; solution-

cast films contain gamma and beta phase; and melt crystallized

films form alpha phase. Kwei’s model was used to model the

glass transition temperatures of the blends. Addition of FCP

increases hydrophobicity of the electrospun membranes. © 2019

Wiley Periodicals, Inc. J. Polym. Sci., Part B: Polym. Phys. 2019, 57,

312–322

KEYWORDS: crystallinity; electrospinning; fluorinated copolymer;

glass transition; PFDMA; PVDF; structure-property relations;

thermal properties; fibers; membranes

INTRODUCTION Separation of oil–water mixtures is a veryimportant technological problem1–4 with many significant appli-cations. Filtration and reuse of wastewater as a drinking watersupply is a major challenge due to its widespread demand5–7

arising from depletion of natural sources and ever-increasingpopulation. From an eco-friendly point of view, it is also the keychallenge in cleaning up oil spills.8,9 Specifically engineeredoleophilic membranes could provide one solution throughselective absorption of oil and rejection of water. The effective-ness of these membranes depends on properties such as highporosity,10 hydrophobicity, oleophilicity,11 and roughness atmultiple length scales.12 Membranes with advanced functionali-ties have been fabricated through the design of functionalpolymers11,13–16 which enhance membrane performance.17–19

Membranes for filtration are manufactured using severalfabrication methods. Common approaches include use of ultra-thin layers of material employing deposition techniques onsubstrates,19–22 spin casting,20 solution-casting,21 and nonsol-vent induced phase separation processes.22–24

Electrospinning (ES) is another popular fabrication methodused to make fibrous membranes,25–28 due to its simplicity andscalability. The approach adopted in this work involves creating

hydrophobic polymer blends and fabricating them into nonwo-ven fibrous membranes by ES. We selected Poly(vinylidenefluoride) (PVDF) as bulk polymer, because PVDF is a verycommon membrane material due to its high chemical andmechanical stability. For the other blend component, wesynthesized a specialty random copolymer, poly(methylmethacrylate-random-1H,1H,2H,2H-perfluorodecyl methacry-late) P(MMA-r-PFDMA) that combines methyl methacrylate(MMA) and PFDMA repeat units along its backbone. MMA waschosen because it is known to be miscible with PVDF,29,30 andwill provide anchoring groups to the membrane matrix. Com-prehensive studies31,32 have shown that compatibility of MMAblended with the amorphous matrix of PVDF is uniform acrossall PVDF concentrations. Therefore, the results of the presentstudy can be extended to copolymers that contain MMA. Thesecond unit, PFDMA, has highly fluorinated side chains thatcould segregate to the ES fiber surface and enhance roughnessand hydrophobicity.33,34 Even though wettability and photo-chemical stability of some copolymers that share a similarstructure to P(MMA-r-PFDMA) have been studied in theliterature,35–37 this specific copolymer is a new addition to thefield. This makes our study unique. Hence, the discoveries of

Additional Supporting Information may be found in the online version of this article.

© 2019 Wiley Periodicals, Inc.

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mailto:Department of Chemical and Biological EngineeringScience and Technology Center4 Colby StreetMedfordMassachusetts02155

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this study are the first findings of blends of this novel copoly-mer, but due to the utilization of MMA in this material, thesefindings can be extended to similar fluorinated copolymers(FCPs) that contain MMA and their blends with PVDF.

Our team has recently reported the copolymer synthesis,performance, and fouling resistance of these membranes.38

However, detailed structure and thermal properties ofPVDF/P(MMA-r-PFDMA) blends have not been reported. Inthis work, nonwoven fibrous membranes were successfullyfabricated from blends of different PVDF/copolymer w/wratios. Our goal is to establish correlations among processinghistory, blend composition, thermal properties, and structure.

Despite the fact that many separate studies link PVDF crystalstructure to its processing history (i.e., by melt crystallization,ES of fibers, or solvent casting), a direct comparison amongthese approaches has not been made. Therefore, we also aimto provide a thorough comparison among these different pro-cessing conditions and the PVDF crystal structures whenblended with the unique FCP blends used in this study.

The morphology of electrospun fiber membranes was deter-mined using scanning electron microscopy (SEM). Crystallo-graphic information was obtained using wide-angle X-rayscattering (WAXS) and Fourier transform infrared (FTIR) spec-troscopy. Thermogravimetry (TG) was used to assess degrada-tion profiles while degree of crystallinity of the blends wasstudied by differential scanning calorimetry (DSC). To assessmembrane wetting properties, sessile drop water contact angletesting was performed on as-spun fibers of all blends, whichbears directly on the industrial relevance of these materials.Solution-cast films were studied for some of the blending ratios,enabling us to draw comparisons to properties of ES fibers.

EXPERIMENTAL

MaterialsPMMA-r-PFDMA was synthesized using free-radical polymeriza-tion by mixing 20 g of total monomer (70/30 w/w ofMMA/FDMA) with 0.02 g azobisisobutyronitrile initiator in 60 mltoluene at 60�C overnight. The reaction was stopped by adding1 g 4-methoxyphenol to the solution and stirring for 15 min. TheFCP was then precipitated in ethanol and washed three times toremove all monomer residues. Then, it was air-dried overnightand further dried in a vacuum oven at 50�C overnight.

The copolymer composition was characterized using protonnuclear magnetic resonance (1H-NMR) spectroscopy38 using aBruker Avance III 500 spectrometer (Bruker Optics Inc., Biller-ica, MA). The copolymer composition was found to be 68/32w/w MMA/FDMA, which was quite close to the nominal compo-sition. Pellets of PVDF (KYNAR® grade 740) were obtained fromArkema Chemicals Inc. (Colombes, France). Blends of PVDF andthe copolymer, named PVDF/FCP, were made by dissolvingthem in common solvents N,N-dimethylacetamide and acetone =1/1 v/v. Nominal composition of the PVDF/FCP blends was con-trolled by varying the PVDF wt % over the following range:100, 95, 90, 85, 80, 75, 50, 25, and 0.

ElectrospinningNonwoven fiber membranes were obtained by ES the solu-tions which had a total polymer content (solids content) of20% w/v at room temperature. The grounded counterelectrode was a flat plate placed 15 cm away, with surfaceperpendicular to the needle, rotating at 1 rpm which resultedin isotropic distribution of the fibers lying predominantly inthe plane of the collector. Electrospun fibers were depositedfrom solution, through a glass syringe of inner diameter14.6 mm with 18-gauge stainless steel needles, at a flow rateof 2.0 ml h−1 controlled by a syringe pump (BS-8000 12VDCSyringe Pump; Braintree Scientific, Inc., Braintree, MA). Thehigh voltage power supply (ES30P-5w; Gamma High VoltageResearch Inc., Ormond Beach, FL) provided an acceleratingvoltage of 20 kV.

Thin/Bulk Film ProductionFor comparison, solution-cast polymer films (with thicknesses�50–100 μm), referred to as “bulk” films, were obtained byspreading the PVDF/FCP solution on a glass substrate usingan adjustable gap doctor blade, and allowing the solvent toevaporate in a fume hood at room temperature. All composi-tions except PVDF/FCP 50/50 formed continuous films. The50/50 blend was brittle and cracked during drying.

Scanning Electron MicroscopyThe morphology of electrospun fibers was studied using aZeiss EVO MA10 SEM (Carl Zeiss, Oberkochen, Germany),operating at 5 kV. Samples were first coated with Au–Pd alloyfor 90 s using a Cressington Sputter Coater 108 (CressingtonScientific Instruments, Watford, UK). Application softwarepackage ImageJ was used to analyze SEM images and toobtain statistics on fiber size by obtaining fiber diameter datafrom no fewer than 100 fibers for each composition.

Wide-Angle X-Ray ScatteringOne-dimensional WAXS was performed on polymer blendfilms and fibers in reflection mode using a Philips PW 1830powder diffractometer (Malvern Panalytical B.V., Almelo,The Netherlands), operated at 40 kV and 45 mA with X-raysof wavelength λ = 0.1542 nm (Cu Kα). Samples weremounted on standard aluminum holders and were examinedin θ/2θ reflection mode (for θ the half-scattering angle)using a step scan interval of 0.02� step−1 and a scanningrate of 0.01� s−1 from 2θ = 5�–45�. The scattered intensityfrom the empty Al sample holder was subtracted as back-ground. All crystallinity indices presented here are normal-ized to the PVDF fraction in each blend and thereforeshould be understood as the crystallinity fraction of thePVDF component in blend.

FTIR SpectroscopyAbsorbance spectra were examined using attenuated totalreflectance FTIR spectroscopy on a JASCO FTIR-6200 Spec-trometer (JASCO Instruments, Tokyo, Japan). Spectra wereobtained from 600 to 4000 cm−1 at 4 cm−1 resolution with256 scans coadded. Air background was subtracted from thesample spectra.

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Thermal AnalysisTG was performed on a TA Instruments, Inc. (New Castle, DE)Q500 series thermogravimetric analyzer from 25 to 800�C at10 �C min−1 under 50 ml min−1 N2 gas flow using polymerfiber samples of mass 5–15 mg. DSC experiments were per-formed on a TA Instruments Q100 series DSC equipped with arefrigerated cooling system with N2 protection gas flowing at50 ml min−1. Sections of bulk films were cut such that theyfilled standard aluminum DSC sample pans with uniform con-tact across the entire surface. Electrospun fiber samples wererolled up and placed on DSC sample pans using the “fixed-end” method described in previous work39 to ensure maxi-mum contact and uniform heat flow. Sample mass variedbetween 3 and 10 mg. The three runs calibration method40

(empty–empty, empty–sapphire, empty–sample) was followedin the temperature range −80 to 200�C at 10 �C min−1 toobtain high precision heat capacity data of the blends. Atemperature-modulated DSC program of heating from −80 to200�C at 5 �C min−1, with a temperature modulation ampli-tude of �0.796�C and period of 60 s, was used to study thethermal properties of electrospun fibers. The DSC cell was cal-ibrated for temperature and heat flow accuracy using indiumstandard, and for heat capacity using sapphire standard.

Contact Angle TestingWetting properties of the fiber mats were measured using aRamé-Hart contact angle goniometer (Succasunna, NJ). Piecesfrom as-spun fiber membranes were cut and taped onto glassslides and their water contact angle was measured using thestatic sessile drop method at room temperature using a 4 μlwater droplet. The final contact angle was calculated by aver-aging over at least 10 individual contact angle measurementstaken for each blend composition.

RESULTS AND DISCUSSION

Fiber PropertiesFigure 1(a–i) shows the SEM images of electrospun fibers for allthe compositions. Uniform fibers free of defects, such as beads,were obtained for all compositions. As shown in Figure 1(j), theaverage fiber diameters were in the range of 0.4–1.9 μm. Thediameter distributions for all individual compositions showedGaussian or skewed Gaussian distributions (Fig. S1, SupportingInformation). The error bars in Figure 1(j) represent the standarddeviations of each data set. Fibers containing 80% or more PVDF(i.e., copolymer content <20%) exhibited average diameterssmaller than 1 μm while fibers with PVDF <80% had averagediameters larger than 1 μm. The fibers tended to be straighter inthe compositions dominated by PVDF [Fig. 1(a–d)] whereas inthe compositions with a high fraction of copolymer [Fig. 1(e–h)],the fibers adopted a more curled shape. Fibers containing PVDF≤75% [Fig. 1(g–i)] were brittle, and could not be bent or rolledwithout fracturing (e.g., when making samples for TG or DSC).

Crystal Phase IdentificationResults from X-ray scattering experiments are shown inFigure 2(a,b) for bulk films and electrospun fibers, respectively,of all blend compositions (100, 95, 90, 85, 80, 75, 50, 25, and

0 wt % PVDF). Intensity has been normalized by dividing by thehighest peak in the diffractogram. In the bottom curves ofFigure 2(a,b), the pure copolymer (PVDF/FCP 0/100) onlyshows amorphous halos at 2θ = 14.5� and 30� in either films[Fig. 2(a)] or fibers [Fig. 2(b)], indicating it is completelynoncrystalline. These two halos arise from the copolymer con-stituents, MMA,41 and PFDMA and suggest that some phaseseparation has occurred.

In the case of films [Fig. 2(a)], homopolymer PVDF (topmostcurve) displays crystalline reflections observed at 2θ values of18.4�, 20.0� , 26.7� , 36.1�, and 39.0� corresponding to d-spacings of 0.48, 0.44, 0.33, 0.249, and 0.231 nm, respec-tively.42,43 These reflections are characteristic of overlappingpeaks of the alpha and gamma crystallographic phases, andrefer to α(020)/γ(020) at 18.3�, α(110)/γ(110) at 19.9�,α(021)/γ(022) at 26.7�, α(200)/γ(200) at 36.2�, and α(002)/γ(004) at 39.0� . The overlapping X-ray reflections do notallow us to determine whether alpha or gamma phase is dom-inant with the same degree of accuracy as FTIR (presented inthe next section). FTIR provides unique and separate absor-bance peaks44 corresponding to alpha, beta, and gamma crys-tallographic phases that enable us to determine the presenceof these phases with certainty. As the copolymer content inthe blends increases, the crystal peak intensities diminish,while maintaining the same angular positions.

In the electrospun fibers, [Figure 2(b)], a shift in peak positionto higher scattering angle is noted for the major reflection at2θ = 20.55� which is characteristic of PVDF β(110) overlap-ping with γ(021). On the lower angle side of this reflection,the peak attributable to α(020)/γ(020) at 18.3� increased instrength for some compositions of the blended fibers com-pared to films [Fig. 2(a)]. The two α/γ overlap peaks seen inthe films at 26.7� and 39.0� are largely absent in the fibers. Apeak at 2θ = 36.4� can be indexed to β(310), and exists inmost fiber compositions. Just as in the films, the crystallinepeaks decrease, and the copolymer amorphous halo increasesas the copolymer content of the fibers increases.

All WAXS diffractograms obtained were deconvoluted usingGaussian curve fitting to yield a crystallinity index, ϕC

WAXS, asin our previous work.45 The relative area under the crystallinereflections over the sampled 2θ range, Ai, allows determina-tion of ϕC

WAXS using:

ϕWAXSc =

Xi

Ai

AT × f PVDFð1Þ

where the numerator refers to the summation over the areasof the i-crystalline reflections, and the denominator AT is thetotal area under all peaks in the diffractogram. In eq 1, crys-tallinity has been normalized to fPVDF, the PVDF fraction in theblends, to obtain the crystallinity index of the PVDF compo-nent in the blends.

In this technique, all crystal and amorphous X-ray reflectionsare assumed to be Gaussian in nature. The Lorentz-correctedscattered intensity, I(q)q2, is fitted as the sum of Gaussiansand a quadratic baseline using:




I qð Þq2 =Xi

Bi expq−qoið Þ22σ2i

!( )+ Cq +Dq2 ð2Þ

where Bi is the amplitude, q0i is the mean q-vector (q = 4πsinθ λ−1, for θ the half-scattering angle, and λ the X-ray wave-length), σi is the standard deviation, and C and D are theq and q2 coefficients, respectively, of the quadratic baseline.

Even though this Gaussian fitting analysis was performed onall WAXS spectra, for the sake of brevity, only two examples

are shown in Figure 3. Parts (a) and (b) show the deconvolu-tion of the X-ray scans of homopolymer PVDF solution-castfilm and blend PVDF/FCP 90/10 as-spun fiber mat, respec-tively. For these particular samples, crystallinity indices werecalculated from eq 1 to be 0.26 for homopolymer PVDF film,and 0.20 for PVDF/FCP 90/10 as-spun fiber mat.

Both film and fiber X-ray data show that PVDF crystallizationis hindered as copolymer content increases, and the decrease

FIGURE 1 (a–i) SEM images confirm the formation of bead free electrospun fibers, spun at 20% w/v solution concentration. The

nominal wt % of PVDF to copolymer, PVDF/FCP, is: (a) 100/0; (b) 95/05; (c) 90/10; (d) 85/15; (e) 80/20; (f) 75/25; (g) 50/50; (h) 25/75; and

(i) 0/100. The scale bars represent 10 μm. (j) The average fiber diameters for all compositions were calculated using ImageJ with data

from at least 100 fibers from each composition. The error bars represent the standard deviations of the data sets. [Color figure can

be viewed at wileyonlinelibrary.com]



http://wileyonlinelibrary.com


is more significant for fibers. The films appear to containeither alpha or gamma phase crystals, while diffraction from

ES fibers is consistent with gamma or beta phase crystals.Next, we examine the infrared spectra of these materials,which is essential for determining the crystal phase.44,46

Figure 4(a–c) shows FTIR spectra obtained for PVDF/PMMA-r-PFDMA solution-cast films, electrospun fibers, and melt crystal-lized films, respectively. The melt crystallized films wereobtained by removing material from the DSC pans after noni-sothermal cooling of molten fibers. This material was includedin the FTIR analysis since melt cooling of PVDF is known toproduce the alpha crystallographic phase,47 and thus aids ininterpreting the results of FTIR on solution-cast films andelectrospun fibers. Pure copolymer, PMMA-r-PFDMA (i.e.,PVDF/FCP 0/100) is shown in dotted lines in the top mostspectra of Figure 4(a–c). Two characteristic peaks are observedat 1146 and 1730 cm−1, corresponding to C-F stretching vibra-tion and O–C=O ester carbonyl stretching vibration,48 respec-tively. The amplitudes of these peaks diminish systematicallyas the fraction of copolymer in the blend is decreased.

FIGURE 2 (a, b) Normalized WAXS intensity versus 2θ for

amorphous copolymer (dashed curve) and its blends with PVDF

(solid curves). (a) Solution-cast films and (b) as-spun fibers. Miller

indices and crystal phases are indicated.63 Curves are vertically

shifted for clarity. (c) Crystallinity index (φCWAXS) versus PVDF

content in the blends of films and fibers calculated from eqs 1 and

2, using deconvolution and Gaussian curve fitting. The error bars

represent repeated trials from different samples with the same

composition. [Color figure can be viewed at wileyonlinelibrary.com]

FIGURE 3 Examples of deconvolution of WAXS spectra,

presented as I(q)q2 versus q, using Gaussian peak fitting in

reciprocal space. (a) Homopolymer PVDF solution-cast film.

(b) PVDF/FCP 90/10 as spun blend fibers. Crystalline

reflections—thin solid black curves; amorphous halos—dashed

black curves; experimental data—blue curve; baseline—green

curve; summation, using eq (2)—red. [Color figure can be

viewed at wileyonlinelibrary.com]






Solution-cast blend films shown in Figure 4(a), do not displayany characteristic absorption features corresponding to α-phasePVDF crystals (e.g., blend films have no absorption peaks ateither 614 cm−1 or 763 cm−1).44,46 These spectra display only

weak characteristic features of the β-phase of PVDF; namely,the characteristic peak at 1275 cm−1 is merely a shoulder com-pared to reference samples with large beta content.44,46 Theydo, however, show a well-defined peak at 1232 cm−1 andincreased intensity at 811 cm−1 consistent with characteristicγ-phase absorption (indicated in the figure with green dashedlines). A complicating factor is that the copolymer has a peak at1239 cm−1 which is very close to γ-PVDF at 1232 cm−1. How-ever, upon careful inspection of the copolymer’s peak, we cansee that its amplitude will be quite small in the blends with75% or greater PVDF content, and therefore overlap of thecopolymer peak in these blends cannot account for the strongpeak at 1232 cm−1. Combining our X-ray and FTIR results, forthe solution-cast blends β-phase characteristics are absent indiffractograms and weak in the absorption spectra. While X-raypeaks can all be indexed to α/γ overlapping reflections, FTIRshows no signature of α-PVDF, and does show characteristicabsorbance of γ-PVDF. Therefore, we can say that solution cast-ing of blends of PVDF with PMMA-r-PFDMA results in filmscontaining predominantly polar phases of PVDF, with largeramounts of gamma, and relatively smaller β-phase content thanis seen in the ES fibers.

In Figure 4(b), absorbance spectra of as-spun fiber mats areshown. These blends have none of the characteristic peaks orspectral regions associated with α-phase PVDF crystals. Fibersdo display the distinctive and characteristic peak of β-phasePVDF at 1275 cm−1 (indicated with a blue dashed line).

The absorbance peak characteristic of γ-PVDF at 1232 cm−1 isreduced in amplitude in the fiber spectra compared to thesolution-cast films. Combining our X-ray and FTIR results forthe electrospun blend fiber mats, β-phase characteristics arestrongly present in both diffractograms and absorption spec-tra. Alpha phase is completely absent from both diffracto-grams and absorbance spectra of as-spun fibers, while gammaphase is present but with reduced signature in fibers com-pared to solution-cast films.

The melt crystallized films are shown in Figure 4(c). Exceptfor pure copolymer, these films all display the characteristicsof α-PVDF, as expected.47 This includes three distinctivereflections of α-phase at 614, 762, and 974 cm−1 marked inthe figure by red dashed lines. In addition, there are twoimportant regions of the spectrum, which are highly signifi-cant for α-phase detection. These include the region from861 to 807 cm−1, which has a relatively featureless downwardslope, and the region from 1227 to 1117 cm−1, which containsa characteristic triplet of peaks. The impact of the copolymeron the absorption spectra is not apparent until its content inthe blend reaches or exceeds 50%. Then, the spectral regionsfrom 861 to 807 cm−1 and from 1227 to 1117 cm−1 acquiremore characteristic features of the copolymer absorption.However, at least in the case of blends containing 75% ormore of PVDF, the absorption spectra can be regarded as hav-ing diagnostic characteristics of α-phase PVDF.

In a manner similar to homopolymer PVDF, the PVDF/copoly-mer blends follow the same trend in which the formation of

FIGURE 4 Normalized FTIR absorbance versus wavenumber for

PVDF/FCP blends at the compositions indicated. The curves are

vertically displaced for clarity. (a) Solution-cast films; (b) as-

spun fibers; and (c) melt crystallized films. Red, blue, and green

dashed vertical lines represent characteristic peak positions for

α, β, and γ crystallographic phases,44,46 respectively. [Color

figure can be viewed at wileyonlinelibrary.com]





electrospun fibers allows PVDF to crystallize preferentiallyinto the piezoelectric β phase (with some γ) while crystallizingfrom the melt results in formation of α-phase exclusively.These results are consistent with previous studies whichinvestigated crystal polymorphism of PVDF under similarprocessing conditions.

Andrew and Clarke49 studied the polar phase crystal formationin electrospun PVDF using homopolymer PVDF dissolved indimethylformamide (DMF). They found that not only is polarbeta phase preferred when PVDF is electrospun, but also thatthe beta phase content in fibers can be increased by ES fromlow viscosity solutions or under a higher applied voltage forhigh viscosity solutions. Yu and Cebe50 electrospun PVDF withnanoclay additives Lucentite® STN and SWN using DMF/ace-tone mixed solvent and found that polar beta phase is pre-ferred in fiber form regardless of the presence of the additive.Some alpha phase content was observed in fibers made usingthe SWN additive which had no organic modifier. Yee et al.51

electrospun PVDF, using DMF/acetone solvent with tetrabuty-lammonium chloride (TBAC) additive. They also observedpolar beta phase crystals formed in electrospun fibers for bothneat PVDF and PVDF with TBAC, and the amount of beta crys-tals was higher when the additive was present. Furthermore,they detected gamma phase crystals in PVDF films producedby spin casting, similar to our solution-cast films in this workwhich also crystallize in the γ crystallographic phase of PVDFwith some presence of β.

This also agrees with the results of Salimi and Yousefi52 whoprepared solution-cast PVDF films. They compared DMAc andcyclohexanone as solvents and found that both solvent sys-tems resulted in crystals containing trans conformers (foundin gamma and beta crystals) and that DMAc promoted moretrans conformers due to its higher dipole moment. Kimet al.53 followed a similar experimental approach to createsolution-cast films using PVDF and several copolymers con-taining PVDF, with (N-methyl-2-pyrrolidone) as the solvent.They also reported the observation of gamma phase PVDFcrystals dominantly present in the films. Ramasundaramet al.54 added an organically modified silicate (OMS) nanocom-posite (Lucentite® STN) additive to PVDF films made withDMAc. They reported the as-cast films to be almost purelygamma phase regardless of the OMS content. Upon meltingand subsequent cooling of these samples, they found slowlycooled melt samples showed alpha crystals predominantly atlower OMS content and the mixture of alpha and gammaphases at medium and higher OMS content. Gregorio andCapitao55 prepared and studied PVDF/PMMA blend films castfrom DMAc solutions, which were taken beyond the melt andthen cooled controllably, resulting in alpha phase PVDF crys-tals. This agrees with our present results of PVDF blendedwith a random copolymer based on PMMA, where we alsoobserve α-PVDF after crystallizing from the melt [Fig. 4(c)].

Thermal PropertiesTG (Fig. 5) indicates that onset of degradation for PVDF isaround 350�C. Addition of the copolymer reduces the

degradation onset to about 270�C. As shown in Figure 5(a),the degradation occurs in several distinct steps that corre-spond to the degradation of components within the blends.The copolymer, shown by the dashed line, has two major deg-radation steps. The PFDMA component of the copolymershows two major degradation steps (see Fig. S3, SupportingInformation), around 300 and 370�C. MMA has a large majordegradation step also around 370�C,56 and a smaller stepclose to 300�C. Homopolymer PVDF, represented by the solidblack line, also shows two degradation steps: major mass lossstep around 470�C and a smaller step around 370�C. Conse-quently, the degradation step seen around 300�C has contri-butions from PFDMA and MMA whereas the step around370�C has three contributions arising from PFDMA, MMA, andPVDF. Because of this complexity, assigning observed TG stepsto these components requires careful accounting of the stepsheights, in terms of the composition. Results are shown inFigure 5(b). The temperatures corresponding to these degra-dation steps remain nearly constant (see Fig. S2, SupportingInformation).

As shown by Figure 5(a), the mass loss at each of these stepsvaries in proportion to the composition of PVDF/copolymer.For compositions in which the blends have a large fraction ofcopolymer, most mass is lost at the first and second steps,that is, at 300 and 370�C.

On the other hand, the third degradation step (around 470�C)accounts for the major portion of degradation in blends domi-nated by PVDF. Figure 5(b) shows the observed mass loss ateach of the main degradation steps, as a function of composi-tion. The expected loss at each of these steps is shown in theinset of the plot and was calculated by consideration ofthe ratio of each component (PVDF, MMA, and PFDMA) in theblends.

Observed data resemble the expected prediction well, showingonly small deviations in a few cases. (The tabulated data usedto generate this figure are available in the Table S1, Support-ing Information.) The difference between the observed andexpected values for mass loss at the different steps couldoccur due to uncertainties in the blend or copolymer’s nomi-nal composition.

Our treatment of TG data, in which different stages (steps) ofdegradation are assigned to the individual components pre-sent in the blend material, allows comparison with the calcu-lated degradation profiles. This provides valuable insight intowhether the presence of other materials affects the degrada-tion profile of specific components. This analysis techniquecan also be used in other multicomponent systems to charac-terize their degradation processes.

Results of apparent specific heat capacity versus temperatureat 10 K min−1 from DSC experiments are summarized inFigure 6(a,b). Figure 6(a) shows the melt crystallized sampleswhich had α-PVDF crystal structure. The data come from thesecond heating of once-melted fiber samples. A general increaseof the glass transition temperature (Tg) with decreasing PVDF




content (increasing copolymer content) is seen in Figure 6(b),which shows a magnified plot of the Tg region, with small verti-cal lines marking the inflection points of the Tg step.

The Tgs of PVDF/FCP blends were measured from the inflec-tion points in the heat flow rate curves and are presented inFigure 6(c), filled circles. For comparison, the glass transitionsof the blends were calculated using two different models. Thefirst was the well-known Fox equation (eq 3) which attemptsto account for Tg mixing in blends57:

1

Tg=w1

Tg1+w2

Tg2ð3Þ

where w1 and w2 are weight fractions of the constituents con-tributing to the glass transition in the blend, and Tg1 and Tg2are their respective individual Tgs (in K). For the individual Tgs,we used 212 K for PVDF58 and 380 K for P(MMA-r-PFDMA).

Crystalline material does not contribute to the glass transitionprocess, and here, only the PVDF can crystallize. Therefore,the amorphous material contributing to the glass transitioncomprises all of the copolymer, as well as the noncrystallinePVDF. The fraction of noncrystalline PVDF is the w1 in eq 3. Inthe crystalline blends, the proportion of copolymer residing inthe amorphous phase may be larger than its nominal contentin the blend, and w2 must be calculated for each blend compo-sition based on the crystallinity. (A full list of calculated dataused in this analysis is given in Table S2, Supporting Informa-tion.) The results of the analysis of Tg using the Fox equationare shown by the red points in Figure 6(c).

The Fox equation does not provide an accurate fit to theexperimental data. It agrees with the general trend at theextremes of composition, and at low PVDF content, but doesnot match the blend data at high PVDF content.

An alternative model given in eq 4, due to Kwei,59 was also usedto calculate the glass transition temperatures of the blends:

Tg =w1Tg1 + kw2Tg2

w1 + kw2+ qw1w2 ð4Þ

In this model, k and q are two fitting parameters that can betaken as quantities characterizing the strength of intermolecu-lar interactions. Here, w1, w2 and Tg1, Tg2 have the samemeanings as before.

The Kwei model, with a single set of fit parameters k and q,gave a better fit to the data over the whole composition rangethan the Fox equation, as shown in Figure S4 (SupportingInformation). Much better agreement to the observed Tgs wasobtained by using two pairs of Kwei parameters. In the regionwith PVDF content 85% and above, k = 0.04 and q = 12 gave agood fit to the experimental data. In the region below 85%PVDF content, k = 1 and q = 107.3 resulted in a satisfactory fit.This suggests that there are relatively weak intermolecularinteractions between the PVDF and copolymer molecules whenthe PVDF is present in large quantity, but relatively strongerinteraction when the PVDF fraction is smaller than 85%.

The Kwei treatment used here is widely employed in thestudy of glass transition temperature and thermal analysis,and can be used to evaluate any mixture comprising miscibleblend partners.

The crystallinity, ϕDSCC , determined from DSC analysis of the

endothermic heat flow, was calculated from eq 5:

ϕDSCC =

ΔHDSCf

f PVDF ×ΔHLitf

ð5Þ

where ΔHDSCf is the measured endotherm area, ΔHLit

f is the lit-

erature value of the heat of fusion of 100% crystalline PVDF(104.6 J g−1)60, and fPVDF is the fraction of PVDF in the blend.Crystallinity values are shown in Figure 6(d) for as-spunfibers (open symbols) and melt crystallized films (solid sym-bols) (see Fig. S5, Supporting Information for the DSC heatcapacity curves for fibers during the first heating scan). As

FIGURE 5 (a) Thermogravimetric analysis on electrospun blend

fibers. Mass remaining versus temperature is shown during

heating at 10 �C min−1 for the compositions indicated.

(b) Observed mass loss at each of the main degradation steps

shown in (a), as a function of composition. Each color

represents the percentage of mass loss of a given component at

a given step. PVDF at 470 �C (green); PVDF at 370 �C (white);

MMA at 370 �C (blue); PFDMA at 370 �C (yellow); MMA at 300 �C(gray); and PFDMA at 300 �C (red). The inset shows the

expected mass loss at the same steps. The steps were assigned

to degradation of each component by comparison to

degradation steps observed in TG curves of homopolymer

PVDF, MMA, and PFDMA (Fig. S3, Supporting Information).

[Color figure can be viewed at wileyonlinelibrary.com]





PVDF content increases, ϕDSCC generally increases for both

electrospun fibers and melt crystallized films. The DSC crystal-linity is larger than the crystallinity index found from WAXS[Fig. 3(c)], which is to be expected. WAXS crystallinity index iscalculated over a limited range of 2θ angles (here, 2θ ≤ 45�),and therefore reflects a lower limit to the overall sample crystal-linity. In addition, WAXS crystallinity index is inherently lessaccurate than DSC-based crystallinity measurements. DSC mea-surements are based on endotherm analysis, which captures themelting of all crystals within the sample. WAXS analysis relieson peak area measurements; only crystallographic planes thatsatisfy Bragg’s Law will contribute to the scattered intensity,again resulting in reduction of calculated crystallinity.

As-spun fibers were crystalline once the PVDF content wasgreater than or equal to 25%. For melt crystallized films, crys-tallization occurred once the PVDF content was greater thanor equal to 50%. It was observed that the crystallinity valuesfor almost all compositions are larger in electrospun fiber

samples than in melt crystallized samples. ES fibers containmostly beta phase crystallographic phase. It has been shownthat fast cooling promotes beta phase crystal growth inPVDF.61 Since the material is being quickly solidified duringthe ES process (within seconds), it is likely that the fast solidi-fication and the high orientation of the fibers promotes a simi-lar effect, where more beta phase crystals are formedresulting in a greater overall crystal fraction.

Crystallization during cooling from the melt occurs around130�C (�400 K). (See Fig. S6, Supporting Information for theDSC cooling heat flow rate curves.) Only blends with >50%PVDF crystallized during cooling. A general decrease of boththe melt crystallization temperature and the exotherm area isseen as the PVDF content in the blends is reduced. Upon sub-sequent reheating, peaks corresponding to melting of thesecrystals [Fig. 6(a)] were observed around 170�C (~440 K).

Crystallization during heating (the so-called “cold” crystalliza-tion) occurred only for PVDF/FCP 50/50, at about 100�C

FIGURE 6 (a) Apparent total specific heat capacity versus temperature during heating at 10 K min−1. Data reflect the second heating

following controlled cooling from the melt at the same rate. Peaks at ~440 K correspond to the alpha phase crystal melting. The

curves are displaced vertically for clarity. (b) Blow-up of the glass transition temperature (Tg) region for clarity. Tg inflection points

are marked with small vertical lines. (c) Observed Tg values (blue circles) from the scans shown in (b) compared to theoretical

predictions for Tg calculated using either the Fox equation57 (red diamonds) or Kwei’s model59 (dashed and continuous curves).

(d) Crystalline fraction of the PVDF component versus PVDF content in blends, determined from the area of the melting endotherm

using eq 3. As-spun fibers (open symbols); melt crystallized material (filled symbols) from the second heating scans shown in (a).

[Color figure can be viewed at wileyonlinelibrary.com]





(~370 K). This suggests that during cooling at 10 K min−1,small nuclei may have formed, but they were not able to growlarge enough to produce a perceptible exotherm signal possi-bly because of proximity of the crystallization temperature tothe glass transition temperature, which is 340 K in this blend.Upon subsequent reheating, these nuclei continued to crystal-lize above the glass transition temperature.

Average water contact angles for as-spun fibers as a functionof the PVDF content in the blends are shown in Figure 7. Withthe incorporation of the highly FCP, the water contact angle isseen to increase significantly, compared to homopolymerPVDF, indicating increased hydrophobicity. This is due to theincorporation of long fluorinated side groups of PFDMA,which segregate to the membrane surface and providing thedesired rough surface chemistry.62 Contact angle generallyincreases with decreasing PVDF content, reaching a maximumat 25% PVDF, and decreasing in the material with 0% PVDF.

The fiber membranes composed of 0% PVDF (100% FCP) aremechanically brittle and the surface of the fibers is not as“fluffy” as fibers from other compositions. This surface effectmay be the reason for the observed contact angle decrease,despite the high level of fluorination. As presented in ourrecent complementary publication38 on the oil–water separa-tion and fouling resistant performance of these membranes,contact angle tests performed with oil droplets show that themembranes modified with FCP exhibit superoleophilicity withoil contact angles <1� .

CONCLUSIONS

We have successfully shown that the proposed blend mate-rials are suitable for production of electrospun nonwovenmembranes. We have investigated the thermal behavior ofblends of PVDF with a random FCP, including the study ofthermal degradation, glass transition temperature, crystal for-mation, crystal polymorphism, melting behavior, and wetting

properties. The onset temperature of degradation decreasesas the copolymer is added to PVDF, and each of the thermaldegradation stages were successfully assigned to the differentcomponents in the blends, in agreement with theoreticalcalculations.

Crystal melting temperatures and the crystal fraction of thePVDF component both decrease, whereas the apparent Tgincreases, as the copolymer content of the blends increases.By fitting the variation of the Tg using the Kwei model, weconclude that there are relatively weak intermolecularinteractions between PVDF and FCP molecules when PVDF ispresent in large quantities; relatively stronger interactionoccurs when the PVDF fraction is smaller than 85%.

Examination of FTIR and X-ray spectra shows that electrospunfibers contain predominantly beta phase PVDF crystals. Meltcrystallization promotes alpha phase crystal growth, andsolution-cast films contain mostly gamma and some betaphase crystals.

Contact angle testing shows that the addition of FCP increasesthe water contact angle of the fibers, indicating greater hydro-phobicity. In a complementary publication, the performance ofthese electrospun fiber membranes in oil–water separationapplications has been reported to be excellent.38 These twopublications report the complete first findings on the synthe-sis, structure and physical properties, and excellent perfor-mance of this novel, highly FCP and its blends with PVDF. Themethods used to investigate these new materials can bedirectly extended to study materials with similar structuresand/or properties, especially polymer blends that containboth semicrystalline and amorphous components. Completeunderstanding of the crystalline phase structure and thermalproperties is necessary to the utilization of electrospun mem-branes for targeted applications.

ACKNOWLEDGMENTS

Support for this research was provided by Tufts Universitythrough a Tufts Collaborates grant; National Science Foundation,Polymers Program of the Division of Materials Research, underDMR-1608125; and the MRI Program under DMR-0520655which provided thermal analysis instrumentation. The authorsthank David Kaplan from the Department of Biomedical Engi-neering at Tufts University for use of the scanning electronmicroscope and Lorenzo Tozzi for instrument training.

REFERENCES

1 R. K. Gupta, G. J. Dunderdale, M. W. England, A. Hozumi,

J. Mater. Chem. A 2017, 5, 34.

2 Y. Z. Zhu, D. Wang, L. Jiang, J. Jin, NPG Asia Mater. 2014,

6, 11.

3 Y. B. Peng, Z. G. Guo, J. Mater. Chem. A 2016, 4, 15749.

4 C. H. Lee, B. Tiwari, D. Y. Zhang, Y. K. Yap, Environ.: Sci. Nano

2017, 4, 514.

5 R. B. Jackson, S. R. Carpenter, C. N. Dahm, D. M. McKnight,

R. J. Naiman, S. L. Postel, S. W. Running, Ecol. Appl. 2001, 11,

1027.

FIGURE 7 Average water contact angle versus PVDF content in

the blends, for as-spun fibers using the static sessile drop

method. Each data point represents the average of at least

10 individual measurements. The error bars represent the

standard error of the mean for each data set. [Color figure can

be viewed at wileyonlinelibrary.com]





6 W. A. Jury, H. J. Vaux, In Advances in Agronomy;

D. L. Sparks, Ed.; Elsevier Academic Press: San Diego, CA, 2007;

Vol. 95, p. 1.

7 W. J. Cosgrove, D. P. Loucks,Water Resour. Res. 2015, 51, 4823.

8 J. E. Kostka, O. Prakash, W. A. Overholt, S. J. Green, G. Freyer,

A. Canion, J. Delgardio, N. Norton, T. C. Hazen, M. Huettel, Appl.

Environ. Microbiol. 2011, 77, 7962.

9 F. Mapelli, A. Scoma, G. Michoud, F. Aulenta, N. Boon,

S. Borin, N. Kalogerakis, D. Daffonchio, Trends Biotechnol. 2017,

35, 860.

10 J. H. Pan, H. Q. Dou, Z. G. Xiong, C. Xu, J. Z. Ma, X. S. Zhao,

J. Mater. Chem. 2010, 20, 4512.

11 C. Lee, S. Baik, Carbon 2010, 48, 2192.

12 G. Kwon, E. Post, A. Tuteja, MRS Commun. 2015, 5, 475.

13 Z. P. Zhou, X. F. Wu, Mater. Lett. 2015, 160, 423.

14 F. Liu, M. L. Ma, D. L. Zang, Z. X. Gao, C. Y. Wang, Carbohydr.

Polym. 2014, 103, 480.

15 C. H. Xue, P. T. Ji, P. Zhang, Y. R. Li, S. T. Jia, Appl. Surf. Sci.

2013, 284, 464.

16 C. Du, J. D. Wang, Z. F. Chen, D. R. Chen, Appl. Surf. Sci.

2014, 313, 304.

17 A. Asatekin, S. Kang, M. Elimelech, A. M. Mayes, J. Membr.

Sci. 2007, 298, 136.

18 A. Asatekin, A. M. Mayes, Environ. Sci. Technol. 2009, 43,

4487.

19 I. Sadeghi, J. Kronenberg, A. Asatekin, ACS Nano 2018,

12, 95.

20 H. Yang, Z. G. Wang, Q. Q. Lan, Y. Wang, J. Membr. Sci.

2017, 542, 226.

21 N. Maximous, G. Nakhla, W. Wan, K. Wong, J. Membr. Sci.

2009, 341, 67.

22 P. Kaner, E. Rubakh, D. H. Kim, A. Asatekin, J. Membr. Sci.

2017, 533, 141.

23 S. Y. Wang, L. F. Fang, L. Cheng, S. Jeon, N. Kato,

H. Matsuyama, J. Membr. Sci. 2018, 549, 101.

24 H. Susanto, M. Ulbricht, J. Membr. Sci. 2009, 327, 125.

25 R. Gopal, S. Kaur, Z. W. Ma, C. Chan, S. Ramakrishna,

T. Matsuura, J. Membr. Sci. 2006, 281, 581.

26 D. Aussawasathien, C. Teerawattananon, A. Vongachariya,

J. Membr. Sci. 2008, 315, 11.

27 R. Wang, Y. Liu, B. Li, B. S. Hsiao, B. Chu, J. Membr. Sci.

2012, 392, 167.

28 N. Govinna, P. Kaner, D. Ceasar, A. Dhungana, C. Moers, K. Son,

A. Asatekin, P. Cebe, Polym. Int. 2019, 68, 231.

29 J. Mijovic, H. L. Luo, C. D. Han, Polym. Eng. Sci. 1982,

22, 234.

30 D. R. Paul, J. W. Barlow, R. E. Bernstein, D. C. Wahrmund,

Polym. Eng. Sci. 1978, 18, 1225.

31 B. Hahn, J. Wendorff, D. Y. Yoon, Macromolecules 1985,

18, 718.

32 J. H. Wendorff, J. Polym. Sci.: Polym. Lett. Ed. 1980, 18, 439.

33 A. M. Coclite, Y. J. Shi, K. K. Gleason, Adv. Mater. 2012, 24,

4534.

34 A. M. Coclite, Y. J. Shi, K. K. Gleason, Adv. Funct. Mater.

2012, 22, 2167.

35 M. Lazzari, M. Aglietto, V. Castelvetro, O. Chiantore, Chem.

Mat. 2001, 13, 2843.

36 O. Chiantore, M. Lazzari, M. Aglietto, V. Castelvetro,

F. Ciardelli, Polym. Degrad. Stab. 2000, 67, 461.

37 N. Tirelli, O. Ahumada, U. W. Suter, H. Menzel, V. Castelvetro,

Macromol. Chem. Phys. 1998, 199, 2425.

38 I. Sadeghi, N. Govinna, P. Cebe, A. Asatekin, ACS Appl.

Polym. Mater. 2019, forthcoming.

39 Q. Ma, M. Pyda, B. Mao, P. Cebe, Polymer 2013, 54, 2544.

40 G. Höhne, W. Hemminger, H. J. Flammersheim, Differential

Scanning Calorimetry; Springer-Verlag Berlin Heidelberg:

New York, 2003.

41 R. Lovell, G. R. Mitchell, A. H. Windle, Faraday Discuss. 1979,

68, 46.

42 M. A. Bachmann, J. B. Lando, Macromolecules 1981, 14, 40.

43 J. B. Lando, H. G. Olf, A. Peterlin, J. Polym. Sci. Part A-1:

Polym. Chem. 1966, 4, 941.

44 B. S. Ince-Gunduz, R. Alpern, D. Amare, J. Crawford,

B. Dolan, S. Jones, R. Kobylarz, M. Reveley, P. Cebe, Polymer

2010, 51, 1485.

45 J. Buckley, P. Cebe, D. Cherdack, J. Crawford, B. S. Ince,

M. Jenkins, J. Pan, M. Reveley, N. Washington, N. Wolchover,

Polymer 2006, 47, 2411.

46 X. M. Cai, T. P. Lei, D. H. Sun, L. W. Lin, RSC Adv. 2017, 7, 15382.

47 R. Gregorio, N. Nociti, J. Phys. D: Appl. Phys. 1995, 28, 432.

48 G. Socrates, Infrared and Raman Characteristic Group Fre-

quencies: Tables and Charts; Wiley: Chichester,UK, 2001.

49 J. S. Andrew, D. R. Clarke, Langmuir 2008, 24, 670.

50 L. Yu, P. Cebe, Polymer 2009, 50, 2133.

51 W. A. Yee, M. Kotaki, Y. Liu, X. H. Lu, Polymer 2007, 48, 512.

52 A. Salimi, A. A. Yousefi, J. Polym. Sci. Partt. B: Polym. Phys.

2004, 42, 3487.

53 K. M. Kim, W. S. Jeon, N. G. Park, K. S. Ryu, S. H. Chang,

Korean J. Chem. Eng. 2003, 20, 934.

54 S. Ramasundaram, S. Yoon, K. J. Kim, C. Park, J. Polym. Sci.

Part. B: Polym. Phys. 2008, 46, 2173.

55 R. Gregorio, R. C. Capitao, J. Mater. Sci. 2000, 35, 299.

56 G. Cardenas, C. Retamal, L. H. Tagle, Thermochim. Acta

1991, 176, 233.

57 T. Fox, Bull. Amer. Phy. Soc. 1956, 1, 123.

58 K. Loufakis, Advanced Thermal Analysis of Fluorinated and

Chlorinated Polyethylenes, Ph.D Thesis, Renesselaer Polytechnic

Institute, Troy, New York, 1986.

59 T. K. Kwei, J. Polym. Sci.: Polym. Lett. Ed. 1984, 22, 307.

60 G. Steiner, C. Zimmerer, Polymer Solids and Polymer Melts –

Definitions and Physical Properties I. Landolt-Börnstein - Group

VIII Advanced Materials and Technologies; Springer-Verlag Ber-

lin Heidelberg, 2013, http://materials.springer.com/lb/docs/sm_

lbs_978-3-642-32072-9_48; Vol. 6A1 (accessed on Mar 01, 2018).

61 A. Gradys, P. Sajkiewicz, S. Adamovsky, A. Minakov,

C. Schick, Thermochim. Acta 2007, 461, 153.

62 J. A. Mielczarski, E. Mielczarski, G. Galli, A. Morelli,

E. Martinelli, E. Chiellini, Langmuir 2010, 26, 2871.

63 K. W. Choi. MSc. Thesis, Cornell University, Ithaca,

New York, 1981.



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