Formation of Oriented Polar Crystals in Bulk Poly(vinylidene fluoride

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Interface-Rich Materials and Assemblies

Formation of Oriented Polar Crystals in Bulk Poly(vinylidene fluoride)/High-Aspect-Ratio Organoclay Nanocomposites Adam Kiersnowski, Kiriaki Chrissopoulou, Philipp Selter, Dorota Chlebosz, Binyang Hou, Ingo Lieberwirth, Veijo Honkimaki, Markus Mezger, Spiros H Anastasiadis, and Michael Ryan Hansen Langmuir, Just Accepted Manuscript • DOI: 10.1021/acs.langmuir.8b02412 • Publication Date (Web): 12 Oct 2018 Downloaded from http://pubs.acs.org on October 16, 2018

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Formation of Oriented Polar Crystals in Bulk Poly(vinylidene fluoride)/High-Aspect-Ratio Organoclay Nanocomposites Adam Kiersnowski1,2*, Kiriaki Chrissopoulou3, Philipp Selter4, Dorota Chlebosz2, Binyang Hou1,5, Ingo Lieberwirth1, Veijo Honkimäki6, Markus Mezger1, Spiros H. Anastasiadis3,7, and Michael Ryan Hansen1,4* 1. Max Planck Institute for Polymer Research, Ackermannweg 10, D-55128, Mainz, Germany 2. Faculty of Chemistry, Wroclaw University of Science and Technology, Wybrzeze Wyspianskiego 27, 50-370 Wroclaw, Poland 3. Institute of Electronic Structure and Laser, Foundation for Research and Technology-Hellas, P.O. Box 1527, 711 10 Heraklion Crete, Greece 4. Institute of Physical Chemistry, Westfälische Wilhelms-Universität Münster, Corrensstr. 28/30, D-48149 Münster, Germany 5. Department of Chemistry and Physical Science, Mount Vernon Nazarene University, 800 Martinsburg Road, Mount Vernon, Ohio 43050, United States. 6. European Synchrotron Radiation Facility, ESRF 71 avenue des Martyrs, 38000 Grenoble, France 7. Department of Chemistry, University of Crete, P.O. Box 2208, 710 03 Heraklion Crete, Greece *correspondence:

Adam Kiersnowski, [email protected], Michael Ryan Hansen, [email protected]

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Keywords: Poly(vinylidene fluoride), chevron-like crystal morphology, organoclay, X-ray scattering, solid-state nuclear magnetic resonance


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Abstract We have investigated the formation of lamellar crystals of PVDF in the presence of oriented clay particles with different aspect ratios (AR) and surface properties. Hot-melt screw extrusion of PVDF with 5 wt. % of montmorillonite (AR~12) or fluoromica (AR~27) resulted in formation of phase separated blends. Replacing the clays with their organoclay derivatives, organomontmorillonite or organofluoromica, resulted in the corresponding intercalated nanocomposites. The organoclays induced formation of polar - and - polymorphs of PVDF in contrast to the -polymorph, which dominates in the pure PVDF and the PVDF/clay blends. Solid-state NMR revealed that the content of the -phase in the nanocomposites was never higher than 7 % of the total crystalline phase, while the : mass ratios was close to 1:2, irrespectively of the AR or crystallization conditions. X-ray diffraction showed that the oriented particles with larger AR caused orientation of the polar lamellar crystals of PVDF. In the presence of the organofluoromica, PVDF formed a chevron-like lamellar nanostructure, where the polymer chains are extended along the extrusion direction, while the lamellar crystals were slanted from normal to the extrusion direction. Time-resolved X-ray diffraction experiments allowed the identification of the formation mechanism of the chevron-like nanostructure.


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Introduction Electroactive, vinylidene fluoride (VDF) – based semicrystalline homo- and copolymers are attracting attention because of their potential applications in sensors, actuators, transducers, memory storage or energy harvesting materials.1-3 The unique electrical properties of VDF-based polymers originate from the presence of crystals (ordered domains) with permanent dipole moments, referred to as polar crystals. The polarity is a result of a particular conformation of VDF units along the polymer chains packed in ordered domains. VDF-based polymers are known for their rich polymorphism. However, only crystals of polar polymorphs are electroactive. The homopolymer, poly(vinylidene fluoride) (PVDF), may occur in four distinct polymorphs called , ,  and  also referred to as Form II, I, III and IV, respectively (Roman numeration follows the order of the discovery). The three polymorphs -, - and - are polar. The crystal structures and phase transitions of most of PVDF polymorphs were investigated as early as in the 1970s and 1980s and were reviewed and supplemented later in the 1990s.4-8 The main chains of homogeneously nucleated, melt-crystallized PVDF may adopt alternating trans-gauche conformation (TGTG’) packed into a monoclinic (P21/c, C52h) cell with dimensions of a=0.496 nm, b=0.964 nm and c=0.462 nm, and the right -angle.4, 9 This non-polar, paraelectric polymorph is referred to as the -polymorph (-phase, -PVDF). The -polymorph is mostly formed after a conventional melt-processing (e.g. injection-molding or extrusion) of the bulk material. This polymorph is the most common in commodity VDF-based plastics. Heterogeneous nucleation or post-crystallization treatment, like stretching and/or poling at a temperature below the melting point triggers the formation of one of the polar crystalline phases. In the polar crystals, the PVDF chains are either in all-trans conformation (TTTT’, -polymorph) or are trans-dominated with mostly periodically distributed gauche- conformers (TTTGTTTG’, 4 ACS Paragon Plus Environment

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polymorph).10-15 Addition of nanoparticles like clays or organoclays is also known to enhance formation of polar crystals in the bulk VDF-based polymers.13, 16-18 Apart from the aforementioned ways, formation of the polar crystals may also occur during crystallization from solution by controlling the solvent evaporation rate or the specific polymer-solvent interactions.19, 20 In the -polymorph, PVDF chains are packed in a body-centered orthorhombic crystalline unit (Cm2m, C14 2v). The reported unit cell dimensions of the oriented specimens containing determined from X-ray and electron diffraction -crystals are a=0.858 nm, b=0.491 nm, and c=0.256 nm.4, 9 Because of the non-trivial packing of the chains within the unit cell, the crystal structure of the polymorph raised controversies in the literature. It seems, however, that the C2cm orthorhombic cell with dimensions of a=0.497, b=0.966 nm, and c=0.918 nm determined by Weinhold, Litt, and Lando from X-ray fiber diagrams is the most widespread in the literature.6 The polarity of - and crystals of PVDF is a source of their pyro- or (negative) piezoelectricity, which make these polymorphs interesting in a variety of modern applications in which transforming mechanical stimuli into electric charge (or vice versa) is required.1, 21, 22 The conformation and packing of the polymer chains in the remaining polymorph, namely the phase, is similar to that in the -phase. Results of X-ray diffraction and spectroscopy indicate, however, that the orientation of the main chains in the unit cell of -phase is distinctly different than in the -polymorph. The unit cells of the truly orthorhombic -phase are not centrosymmetric, as all the fluorine atoms are “pointing” to the same direction with respect to the backbone zig-zag plane: this gives a permanent dipole moment to the -crystals and makes them polar.23 The structure of the -crystals makes PVDF ferroelectric and hence, interesting for applications in, for instance, non-volatile memories.23


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To take advantage of the electroactive properties of PVDF in for example sensors or transducers, it is necessary not only to obtain a specific polymorph, but also to make the crystalline phase uniformly oriented on a scale corresponding to the active elements of the devices.24, 25 This makes all methods of film or fiber formation suitable in this area since both films and fibers typically reveal a high level of anisotropy. For instance, oriented films of -phase-rich PVDF can be formed as a result of epitaxial crystallization of the polymer on potassium bromide crystals.9 Oriented phase can also be formed as a result of biaxial drawing of free-standing PVDF films formed either from solution or the melt; this method is currently used in practice to produce commercially available piezoelectric PVDF films (foils). Electrospinning of PVDF with addition of organoclays leads to the formation of highly anisotropic fibers with enhanced -phase content as well.26, 27 Nevertheless, having the oriented polar PVDF crystals only in the form of films or fibers, can actually be considered a limitation in practical use. A challenge would be formation of piezoelectric PVDF-based materials of various shapes by large-scale processing techniques, like injectionmolding or 3D printing.28 This could broaden the range of applications of PVDF to encompass, for instance, new areas in robotics or sensorics, (e.g. soft, freely shaped sensoric ‘skin’ for robots) or soft electric generators in smart clothing products. The development of such materials would require finding a controlled and reproducible way to force the formation of anisotropic polarcrystalline PVDF upon solidification after the melt processing of the polymer. In this work, we focus on the large-scale orientation of polar crystals of PVDF achieved by constraining the polymer flow patterns in between flat particles of organically modified fluoromica (organofluoromica). The fluoromica is a synthetic smectite clay mineral with particles having an aspect ratio (average diameter-to-thickness) as high as ~30, which is more than twice as large as the aspect ratio of naturally occurring clay particles.29 It is well known that mixing


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polymers with layered inorganic materials, such as smectite clays, can lead to three different types of structures, depending on the specific interactions between the two components. Of these structures, we can distinguish the phase separated systems, where the nanoparticles remain tightly aggregated in the form of micron-sized clusters. Tuning the interactions in the system may cause intercalation of polymer chains into the galleries of stacked inorganic platelets and hence, integration of structure at the nanoscale.

30, 31

In the case of favorable attractive interactions

between the macromolecules and the surface of the nanoparticles, the latter may disaggregate and disperse in the polymer matrix.32 This is the way how exfoliated clay-polymer nanocomposites are formed.33,

34

Organoclays, resulting from modification of surfaces clay minerals with

amphiphilic molecules, are known for their potential in enhancing mechanical, thermal, or barrier properties as well as reducing the flammability of plastics. Smectite clays are also known to have influence on crystallization behavior of semicrystalline polymers, including the ability to affect their anisotropy, nucleation rate, degree of crystallinity, or even altering their crystal structures.3537

In the context of this work, it is particularly important that organoclays are already known to

generate - or -crystals of PVDF.18, 38 In order to supplement the knowledge in this field, we demonstrate that high-aspect ratio organoclays may additionally promote formation of oriented, lamellar crystals of PVDF. The orientation effects observed here are similar to those reported by Beaucage for HDPE-organoclay films.36 In our approach, however, it is not necessary to use conventional anisotropy-forcing techniques such as extrusion of the materials into the form of sheets, high-speed (electro)spinning or stretching of the solid specimens. Our results prove that simple, large-scale processing techniques, such as screw extrusion as applied here, can also be effective in producing anisotropic, crystalline PVDF-based nanocomposites. Thus, it seems possible that the method described here may underpin technologies for the production of


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anisotropic, PVDF-based materials, enabling applications in fields where electroactive polymer materials are required. Experimental Section Materials Poly(vinylidene fluoride) (PVDF, Mw=180 kg/mol, Mn=71 kg/mol) was purchased from Sigma-Aldrich (product no. 427152) and was used as received. We utilized two different clays: 1) sodium montmorillonite (Nanofil®116, MT) with a cation exchange capacity of 1.16 meq/g that was a gift from the German Division of Southern Clay Products in Moosburg, and 2) sodium fluoromica (SOMASIF® ME100, further referred to as FM) with cation exchange capacity of 1.15 meq/g (CO-OP Chemical Japan).29 The clays (20 g of each) were modified (organophilized) via a conventional ion-exchange reaction using the equimolar (1:1 mol:meq) amounts of dodecyltrimethylammonim bromide (Sigma-Aldrich) as described elsewhere.29 The ionexchanged products were washed with water and ethanol, subsequently dried at 80C in air and after grinding into a fine powder they were further dried in vacuum. The resulting organoclays were abbreviated as OMT and OFM for organomontmorillonite and organofluoromica respectively. The positions of basal X-ray reflections and basal distances of the clays and organoclays are shown in Figures S1 and S2 in the supporting information (SI). Preparation of the composites and extrusion of rods The composites were prepared using the 5 cm3 DSM -processing twin screw micro-extruder (S/N 98001, Figure 1a) according to the following protocol: 5.7 g of PVDF and 0.3 g of either MT, FM, OMT or OFM were added to the extruder heated to the constant temperature of 215C. The ingredients were added and premixed for approx. 150 seconds. Then the systems were blended (recirculated) for additional 300-350 seconds at the rotor speed of 150 rpm followed by 8 ACS Paragon Plus Environment

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extrusion through a 1.7-mm die (Figure 1a). After cooling to room temperature, the extrudates were cut into 30-mm rods (Figure 1b). The (nano)composites were abbreviated as PVDF/MT, PVDF/FM, PVDF/OMT or PVDF/OFM; they were all containing 5 wt. % of filler. X-ray scattering measurements The static X-ray scattering measurements were performed with point collimated, monochromatic copper radiation (8.04 keV, =0.1542 nm) generated either by the sealed 1.2 kW X-ray tube (Siemens) in the wide-angle scattering (WAXS) or the 1.5 kW rotating anode (Rigaku MicroMax-007) in the small-angle scattering experiments (SAXS). The radiation scattered by the samples was recorded on area detectors: MAR345 from Marresearch in WAXS or the HI-STAR multiwire proportional chamber from Bruker in SAXS. In both static WAXS and SAXS experiments the camera length and therefore the q-range (q=4sin-1) were calibrated using the silver behenate standard. The experimental geometry is schematically shown in Figure 1c.


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Figure 1. The micro-extruder in (a) was used to prepare the 301 mm rod-like samples shown in (b). The experimental geometry used in all the X-ray scattering measurements with the rod-like samples aligned parallel to the z-axis (c). The diameter (D) of the extrusion die was 1.7mm, while the channel length (L(c)) = 30mm. The 1 eurocent coin in (b) is included to get an impression of the sample sizes.

Time-resolved wide-angle X-ray scattering experiments were performed using the instrumentation available at the ID15 synchrotron beamline (European Synchrotron Research Facility, Grenoble, France). The beam energy (37.8 keV, =0.033 nm) and the camera length (115.4 cm) were calibrated using the (200) and (111) peaks of cerium dioxide (CeO2) standard. The scattered radiation was recorded on MAR345 detector. The samples were wrapped in aluminum or kapton and mounted in the heated holder enabling transmission of X-rays; the holder temperature was controlled by the Lakeshore 336 with 0.02C resolution. The samples 10 ACS Paragon Plus Environment

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were heated from 40C to 200C at a constant rate of 5C/min, thermally equilibrated for 5 minutes and then cooled to 40C with 1C/min ramp. One-dimensional scattering profiles were obtained by appropriate integration of the two-dimensional patterns. The scattered intensity was plotted as a function of either the scattering vector (q) or the azimuthal angle () (see Figure 1). Electron microscopy imaging For scanning electron microscopy, the unmodified clay particles were deposited from a dilute (10 g/L) water (Milli-Q,18.2 MΩ) dispersion on the gold substrates having a roughness of 0.2 nm. The images were recorded on a Hitachi SU8000 microscope (2.7 kV). For transmission electron microscopy the samples were cut along the z-axis (see Fig. 1 c) into 50-nm ultra-thin sections and imaged using a FEI Tecnai F20 transmission electron microscope (200 kV). Spectroscopic measurements Fourier-Transform Infrared (FTIR) measurements were performed using a Bruker Vertex 70v infrared spectrometer fitted with horizontal ATR device (Ge, 45°). For evaluation, 64 scans with 4 cm–1 resolution in the mid-wavelength infrared (MIR, 4000-400 cm–1) and 16 scans with 2 cm−1 resolution in the near-infrared (NIR, 600-30 cm–1) range were taken. The specimens for the

19F

solid-state NMR measurements were prepared by melting the

samples (5-10 mg) at 200°C followed by cooling (-50°C/min) to and annealing (3 hours) at two crystallization temperatures in the Mettler-Toledo DSC 1 system. The selected crystallization temperatures (see Tab. 1) correspond to the shortest and the longest observable isothermal crystallization times. After annealing, the specimens were cut into ~0.1 mm pieces and then packed into ZrO2 rotors with outer diameter of 2.5 mm. All

19F

solid-state NMR spectra were

acquired on a Bruker AVANCE III 500 spectrometer operating at 11.7 T, corresponding to a Larmor frequency of 470.6 MHz for 19F. A commercial 2.5 mm H/F/X MAS probe employing a 11 ACS Paragon Plus Environment

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spinning frequency of 25.0 kHz was used. Radio-frequency field strengths and

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19F

chemical

shifts were referenced to PTFE (-122 ppm) as an external reference. The magic angle was set using KBr before each measurement.39 The

19F

/2 pulse length was 2.5 µs and the spin-lock

(SL) pulse length was set to 20.0 ms. High-power 1H decoupling using the swfTPPM scheme40 was employed at a radio-frequency field strength of 100 kHz. A recycle delay of 10 s was used for all samples. Data processing was performed using the Bruker Topspin software package and spectral deconvolution was performed using dmFit.41 For the deconvolution, the 19F MAS NMR spectrum obtained using a spin-lock pulse of 20.0 ms for PVDF/OFM, 158°C was chosen for the initial fit, since this sample contained all three polymorphs in sufficient quantities to accurately determine the peak positions. Seven Gauss/Lorentz (G/L) peaks were employed for the three different polymorphs (one pure α peak, one peak containing at least both α and β, as well as possibly γ contributions, three peaks for the γ polymorph, and two further peaks accounting for the broad should around ~-100 ppm, with possible γ contributions, vide infra). and another broad G/L peak was used for modeling the polymer defects (region-irregular polymer domains, corresponding to head-head/tail-tail defects42). All peak positions and relative intensities were based on the values previously reported by Koseki et al. as well as Hucher et al..43, 44 All fits used the same starting conditions obtained from the initial fit, peak positions were allowed to vary slightly (ca. 0.5 ppm). The G/L ratios were fixed to 0.5 for all signals except for the polymer defects, where a larger Gaussian contribution was assumed, i.e., the G/L ratio was fixed to a higher value of 0.9. Peak assignment to the individual phases was based on literature values. However, ambiguities in the assignment of the broad signal at ~-100 ppm lead to the development of two models (vide infra).


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Results and discussion In order to summarize all observations and findings underpinning the final conclusions, we will in the following three sections discuss: i) the morphology of the investigated clay particles, ii) the influence of clay and organoclay particles on the structure and anisotropy of PVDF crystals, and iii) the influence of clay or organoclay particles on crystal polymorphism of PVDF in the (organo)clay-containing blends or nanocomposites. The experiments were designed to provide insights into the phase composition of the materials and also to find an explanation for the formation mechanism of the oriented PVDF crystals. Morphology of the clay particles


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Figure 2. Scanning electron micrographs of (a) MT and (b) FM particles on gold surface and transmission electron micrographs of isolated (c) MT and (d) FM particles in the PVDF matrix. Cartesian axes in (c) and (d) are in accordance with Figure 1c.

Analysis of electron micrographs (Figure 2a and 2b) indicates that the montmorillonite (MT) particles mostly have diameter around 1 m while the diameter of the fluoromica (FM) particles is larger, reaching roughly 3 m. After blending with PVDF and extrusion of the rods (Figure 1b), the average diameters of both MT and FM particles noticeably drop yet still differ by a factor of approx. 3 (Figure 2c and 2d). The aspect ratios of the clay particles, approximated on the basis of TEM images were found to be 12 and 27 for MT and FM respectively (Figure 2c and 2d). According to the additional TEM studies (see Figure S3 in SI), the shearing forces in the extruder were high enough to separate larger particles into stacks containing ~10-20 clay platelets and disperse them in the polymer matrix. In accordance with TEM, the results from WAXS measurements (Figure 3) of either PVDF/MT or PVDF/FM systems suggest that the clay particles remain in the form of stacks. Such a conclusion can be drawn after inspection of the WAXS patterns for both blends; the weak reflections at q~6.2 nm-1 (PVDF/MT) as well as q~5 nm-1 and q~6.5 nm-1 (PVDF/FM), corresponding to basal spacings (dI) of the clays are still visible (insets in Figure 3a and 3b, see also Figures S5 and S6 in SI). FM in the PVDF/FM system reveals a double basal reflection resulting from partial adsorption of water in the interlayer spaces. This effect was discussed in detail in our earlier articles.29, 45

X-ray scattering studies of the poly(vinylidene fluoride)/clay blends and poly(vinylidene fluoride)/organoclay nanocomposites The WAXS profile shown in Figure 3a suggests that the addition of MT has no significant influence on the polymorphism of PVDF – the scattering pattern of PVDF is identical to those 14 ACS Paragon Plus Environment

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published in the literature and verified with spectroscopy to be the -polymorph (see Figures S4S6 and Table S1 in SI). Thus, based on WAXS, it can be concluded that similar to the pure PVDF, the polymer in PVDF/MT blend crystallizes mainly in the -polymorph. The practically constant intensities of all Debye-Scherrer rings over the -angle visible in the WAXS pattern (Figure 3a) suggest that the crystals are randomly oriented in the sample volume, and that the system is isotropic on a macroscopic scale. The isotropic orientation and the type of PVDF crystals in the PVDF/MT blend can also be concluded from the distribution of intensity observed at small angles i.e. in SAXS experiments. In the SAXS pattern, shown in Figure 3c, a characteristic halo-like ring with maximum at qmax = 0.51 nm-1 can be seen. This ring results from the Bragg diffraction of X-rays by lamellar crystals alternately stacked with amorphous layers forming remarkable layered structures typical of thermoplastic polymers. The amorphous-tocrystalline distance of the PVDF lamellar structure in PVDF/MT blend (dac, the long period) calculated from the position of peak intensity (qmax) is equal to 12.7 nm. The ring related to the long period reveals the almost constant intensity over the whole -angle range (Figure 3c, inset), which supports the conclusion about the formation of isotopically distributed lamellar crystals in PVDF/MT system. The subtle increase in the intensity observed at the equator (i.e. around =0 and 180, see Figure 3c, inset) is a ‘tail’ of the intensity scattered by the clay particles preferentially oriented along the rod axis (z-axis). As evidenced by the WAXS data (Figure 3b, inset), the extrusion of PVDF/FM blends also causes the formation of the -polymorph of PVDF. In this case, however, the crystal structure reveals a notable degree of anisotropy. In the SAXS pattern of this blend we can see that the intensity scattered by the lamellar structure of PVDF reveals broad, discrete maxima located at the meridian (=90 and 270, Figure 3d). Integration of the intensity over the q-range 15 ACS Paragon Plus Environment

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corresponding to the long period peak to obtain the I() function bolsters this observation. The I(), plotted in the -range of 20-160° (Figure 3d, inset), reveals a maximum at =90°. This indicates that the lamellar crystals of PVDF in the PVDF/FM blend are oriented mainly in the x-y plane. At the same time, the intensity diffusely scattered by the clay, which is visible in the center of the SAXS pattern as a remarkable equatorial ‘streak’, suggests a preferential orientation of FM particles along the z-axis. On this basis it can be concluded that the PVDF lamellar crystals are oriented perpendicularly to the surfaces of clay particles, being mostly aligned with the long axis of the extruded rods. Let us note here, that the intensity diffusely scattered by the FM particles (the ‘streak’) is narrower over the qz axis in comparison to the streak visible in the pattern of PVDF/MT sample (cf. Figure 3c and 3d). This may suggest a sharper axial alignment of FM particles along the z-axis than the one observed in the case of MT particles. We attribute this sharper orientation of the FM particles to their higher aspect ratios and hence their higher tendency to orient along the flow of polymer melt caused by extrusion.


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Figure 3. WAXS and SAXS patterns of (a, c) PVDF/MT and (b, d) PVDF/FM samples. The labels dI and dac stand for interlayer distance of the clay particles and long period of lamellar structure of PVDF, respectively. The insets in (a) and (b) show the I(q) scattering patterns resulting from the integration of the 2D patterns over the whole azimuthal angle () in the q range of 3-25 nm-1. The insets in (c) and (d) show the azimuthal intensity distributions I() within the q-range of 0.4-0.6 nm-1 The Cartesian axes at the center of the Figure illustrate the sample geometry in accordance with Figure 1c.


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Figure 4. WAXS and SAXS patterns of (a, c) PVDF/OMT and (b, d) PVDF/OFM extruded rods. The insets in (a) and (b) show the I(q) scattering patterns resulting from integration of the 2D patterns over the whole azimuthal angle () in the q range of 4.5-19.5 nm-1. The insets in (c) and (d) show azimuthal intensity distributions I() within the qrange of 0.4-0.6 nm-1 The Cartesian axes at the center of the Figure illustrate the sample geometry in accordance with Figure 1c.

The ion-exchange of either MT or FM clays with equimolar (with respect to the cation exchange capacity) amounts of dodecylammonium bromide resulted in the formation of organoclays. The interlayer distances in the organoclays as determined by X-ray diffraction were 1.6 nm and 1.9 nm for the organomontmorillonite (OMT) and organofluoromica (OFM) 18 ACS Paragon Plus Environment

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respectively (see Fig. S1 and S2). Extrusion of PVDF with either OFM or OMT resulted in intercalated nanocomposites since the basal spacing (dI) of the organoclays increased to 2.8 nm (PVDF/OMT) and 2.7 nm (PVDF/OFM) (Figure 4, see also Figure S1, S2, S7 and S8 in SI) from 1.6 nm and 1.9 nm, respectively (Figure 3). WAXS patterns of both PVDF/OMT and PVDF/OFM nanocomposites (Figure 4a and 4b, respectively) indicate that the particles of both the organoclays are sharply oriented along the extrusion direction. The diffraction peaks of the organoclays related to their dI are located at the equatorial positions at q=2.24 nm-1 (PVDF/OMT) and 2.33 nm-1 (PVDF/OFM). Unlike in the blends with the unmodified clays, in both PVDF/OMT and PVDF/OFM nanocomposites the organoclay particles are sharply aligned along the extrusion axis (cf. e.g. Figure 3b and 4b). While the real reason of a difference between orientation of particles in the blends and nanocomposites cannot be explained on the basis of our results, we can hypothesize that aspect ratios and surface properties of the particles may play a major role here. In the blends with the unmodified clays, where no attractive polymer-particles interactions occur, the orientational structure of the clay particles is likely driven mainly by aspect ratio of the particles. Probably the particles of fluoromica with the larger aspect ratio get oriented easier than the particles of montmorillonite. In the case of the nanocomposites, it seems likely that shearing forces upon extrusion cause a usual orientation of the polymer chains, which, after adsorption on the surfaces of the organoclays cause the latter to orient in the direction parallel to polymer chains, i.e. along the extrusion axis. Hence, it seems that attractive interactions between the clay and the polymer are the key prerequisite for a sharp alignment of the organoclay particles along the extrusion flow. Comparing the WAXS profiles shown in Figure 3a, 3b with those in Figure 4a and 4b, shows that the crystal structure of PVDF in the nanocomposites containing either OMT or OFM differs 19 ACS Paragon Plus Environment

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from that of the pure PVDF and that of the blends with the unmodified clays. In addition, the results of DSC measurements of the samples indicate, that the organoclays in either of the nanocomposites cause an increase in the melting points and a decrease in overall crystallinity of PVDF (see Figure S9, Table S2 and also additional comments in SI). The observed changes in thermal properties additionally support the WAXS-based conclusion about a change in the crystallization of PVDF in the nanocomposites. One-dimensional WAXS I(q) profiles (insets in Figure 4a and 4b) recorded for both PVDF/OMT and PVDF/OFM nanocomposites reveal distinct maxima at q=13.2 nm-1 and q=14.5 nm-1 corresponding to reflections of either the - or -phase, or both (see Figure S7 and S8 in the SI).17 Since the analysis of the phase composition based on WAXS results is rather ineffective due to the limited number of independent, non-overlapped diffraction peaks, the quantitative analysis was performed using spectroscopic techniques in accordance with the preceding literature. These results are discussed in the next section. Let us now focus on the morphological features of PVDF crystals in the blends and nanocomposites. Our WAXS and SAXS studies indicate the formation of isotropically distributed polymer crystals in the PVDF/MT blend (Figure 3a) and a certain degree of orientation of crystals in PVDF/FM blend (Figure 3b). Based on both WAXS and SAXS patterns of PVDF/OMT, shown in Figure 4a and 4c, respectively, we conclude that the lamellar crystals of the polymer in the PVDF/OMT nanocomposite are, to some extent, spatially oriented. The intensity of the main diffraction ring in the WAXS pattern is not uniform, but it does reveal an indistinct maximum at the equatorial plane and a minimum on the meridian. In the SAXS pattern, in turn, the halo-like ring reveals a weak maximum at =90 (Figure 4c). Both WAXS and SAXS patterns (Figure 4a and c) indicate that the PVDF lamellae are oriented in the x-y plane, i.e., the lamellae are perpendicular to the fiber axis and surface of the clay particles. The long period of


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the lamellar structure (dac) for PVDF in the PVDF/OMT nanocomposite is 11.2 nm, which is shorter by 1.5 nm compared to dac determined for the PVDF/MT system. The long periods in the systems based on FM or OFM are similarly different: in the PVDF/FM and PVDF/OFM systems we observed a dac of 12.8 nm and 11.4 nm respectively. Considering the differences in the phase compositions with the different additives, one can conclude that lamellar crystals of polar polymorphs (, ) form lamellar stacks with the long period smaller than that observed in the case of non-polar polymorphs (). The intensity distribution in the WAXS pattern of PVDF/OFM (Figure 4b) indicates that the PVDF main chains are parallel to the z-axis of the extruded rods (see Figure 1c). Furthermore, the SAXS pattern (Figure 4d) reveals a pair of maxima symmetrically distributed along the meridian, forming a so-called four-point pattern (Figure 4d). The reflections are shifted away from the meridian (i.e. from =90 by an angle of ±22° (Fig. 4c, inset) As suggested in previous studies, such a remarkable intensity distribution may result from a tilt of the lamellar crystals with respect to the z-axis of the extruded rods, corresponding to the formation of chevron-type arrangements of lamellar crystals.46, 47 Combining WAXS and SAXS results from Figure 4 enables us to propose a complex model for the nanostructure of PVDF/OFM nanocomposite as schematically shown in Figure 5a. A two-dimensional fast Fourier transform (Figure 5c) of the picture shown in Figure 5a reveals an agreement with experimental SAXS pattern (Figure 4d), indicating the validity of the proposed model of the lamellar nanostructure. Formation of such remarkable structures with ‘tilted lamellae’ has previously been reported for liquid-crystalline and semi crystalline systems, including solution-cast poly(Llactide) (PLLA)48, polyurethane,49 and poly(ether-ester) elastomers50 under strain/relaxation conditions. According to Murthy and Grubb46 such morphologies may result from deformation, interlamellar shear at the microdomain level, and chain slippage within the lamellar crystals.


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Figure 5. Schematic drawing of (a) the chevron-like structure for the PVDF/OFM nanocomposite. (b) Structure of the crystalline PVDF lamellae and their orientation at the organoclay interface (black lines represent PVDF chains). (c) Two-dimensional Fourier transform of the model proposed in the figure (a). In (a) the crystalline and amorphous areas are colored dark and light green, respectively. The Cartesian axes in the left corner of the Figure illustrate the sample geometry in accordance with Figure 1c. Note, that the Qx,y in the (c) denotes the dimension of the simulated reciprocal lattice and not a scattering vector distinguished by a lowercase q.

Shear flow applied during the extrusion or spinning of polymers is known to cause linear polymer chains to orient along the shearing direction.27 In the case of semi crystalline polymers, the orientation of the polymer chains in the melt may also contribute to a distinct orientation of the lamellar crystals, including their chevron-type arrangement, as observed here for the PVDF/OFM nanocomposite. Shear flow may also cause orientation of nanoparticles such as nanotubes or nanosheets.37 A strong orientation of OFM or OMT particles along the extrusion axis caused by flow can be concluded from the results of our X-ray diffraction experiments 22 ACS Paragon Plus Environment

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(Figure 3 and 4). Considering the characteristic orientation of organoclay particles, the question arises how they contribute to the orientation of the polymer crystals? This question is particularly justified here in view of the fact that an alternative explanation for a preferred mutual orientation of different crystals accompanied by a change in their crystal structures could be epitaxy. Epitaxy, for instance, is the reason for the formation of heterogeneous shish-kebab crystals formed upon crystallization from solution as well as highly oriented crystalline structures at interfaces, resulting from crystallization of polymer melts.37, 51 The rationale to hypothesize the occurrence of epitaxy in the case of heterogeneous PVDF systems is the formation of the polar  and  crystals, suggesting a different nucleation mechanism than what is found in pure PVDF.52 To verify whether the orientation observed in our experiments results from (i) the orientation of the polymer melt or (ii) the epitaxial growth of crystals nucleated by highly oriented particles, we have performed additional, time-resolved scattering experiments with synchrotron radiation. The PVDF/OFM rod was sealed in a kapton tube and placed in the X-ray beam as shown in Figure 1c. The sample was heated to 200°C (i.e. above the melting point of PVDF) and then slowly (1°C/min) cooled to ensure quiescent crystallization of the polymer. During the melting and crystallization, we have recorded a series of WAXS patterns. The exemplary WAXS patterns recorded at 25°C, 200°C and again at 25°C after cooling and recrystallization, are shown in Figure 6a-c. These WAXS patterns correspond to the sample before melting (Figure 6a), in its molten state (Figure 6b), and recrystallized after melting (Figure 6c). One-dimensional diffraction data, obtained by integrating the 2D WAXS patterns, are plotted as false-color intensity maps either against the azimuthal angle (, Figure 6d) or scattering vector (q, Figure 6e). They are shown in Figure 6 to demonstrate changes in anisotropy and crystal structure over the heating and cooling cycles.


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Figure 6. 2D WAXS patterns of (a) as-extruded, (b) molten and (c) recrystallized PVDF in the PVDF/OFM rod. Xray diffraction intensity shown as 2D false-color maps recorded on subsequent heating and cooling as a function of temperature and (d) azimuthal angle or (e) scattering vector q. The intensities (I’()) in (d) were calculated by dividing the measured intensity at a given  by the intensity measured at the equator (=90°, cf. Figure 1c).

From the scattering data in Figure 6, it is clear that heating the PVDF/OFM nanocomposite to 200°C and beyond causes a gradual deterioration of the anisotropy. Already at 80 °C, we observe a gradual decrease in the intensity of 110 reflections for PVDF (visible in Figure 6d at =306° and 54°), which at 160°C disappears, showing that PVDF within the nanocomposite goes into the isotropic melt. The one-dimensional WAXS patterns recorded before and after the melting (Figure 6e) suggest that recrystallization of PVDF in the PVDF/OFM nanocomposite leads to the identical crystal system regardless of their orientation (Figure 6d). The observed changes in the azimuthal distribution of scattered intensity (Figure 6a-d) further shows that melting and recrystallization from an isotropic melt causes formation of an isotropic system of polymer crystals, irrespectively of the orientation of OFM particles. Orientation of the latter is 24 ACS Paragon Plus Environment

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Langmuir

independent of whether the polymer is molten or solid. The indistinct change in the profile of the basal reflection of OFM is likely related to reorganization of the polymer intercalated in the OFM structure, similarly to what has been reported in earlier work.37 In summary, the above scattering experiments following the melting and recrystallization of the PVDF/OFM sample indicate that the formation of anisotropic PVDF crystals in the nanocomposite is a result of non-static crystallization conditions upon extrusion. On the other hand, the direct interactions of PVDF with the surface of the organoclay is the reason why polar crystals are formed, possibly due to the heterogeneous nucleation.

Spectroscopic studies of the phase composition for the poly(vinylidene fluoride)/clay blends and poly(vinylidene fluoride)/organoclay nanocomposites Distinguishing the -phase from polar - and - polymorphs is relatively easy and can be done using WAXS, because the diffraction pattern of the -phase is clearly different from the other two polymorphs. However, differentiating between the - and - polymorphs using X-ray diffraction is not possible due to the similarity of their WAXS profiles.10, 53 Therefore, the phase composition of PVDF is typically studied by either FTIR or NMR spectroscopies.54 In our approach, PVDF and its composites were first analyzed using FTIR in the attenuated total reflection (ATR) mode. Since the extruded rods were too thick to irradiate through them, they were intersected in the middle and the pieces were pressed onto the germanium window of the spectrometer to ensure a good contact between the window surface and the specimen. The spectra of the pure PVDF sample, shown in Figure 7, recorded both in the medium (MIR, Fig. 7a) and near infrared (NIR, Fig. 7b) reveal bands at 1425, 1382, 1210, 1148, 975, 795, 763, and 614 cm-1 (MIR) and 532, 490, 410, 354 cm-1 (NIR), characteristic for the -polymorph of PVDF


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formed upon extrusion and crystallization from the melt. The lack of any additional IR bands assignable to other polymorphs suggests that PVDF without any additives during extrusion practically results in a phase-pure -phase of PVDF. The addition of MT or FM clay particles has only little influence on the phase composition of PVDF: from analysis of the NIR spectra (Figure 7b) one can conclude that the -polymorph is dominating. However, in the MIR spectra in Figure 7a additional and relatively weak bands are visible at approx. 835 cm-1 for both PVDF/FM and PVDF/MT systems. Based on literature54, this band is considered to be the characteristic vibrations of the electroactive, trans-dominated phase of PVDF, suggesting that the addition of either clay triggers the formation of small amounts of polar PVDF crystals in addition to the dominating -polymorph. Based on the intensities observed for the band at 835 cm-1 one could expect some traces of polar crystals in the WAXS patterns. Indeed, in the one-dimensional WAXS profile of PVDF/MT system (Figure 3a, inset), we can see an additional shoulder at q~15 nm-1 (marked by an asterisk), which can be connected with the presence of either - or -phase in this blend. This shoulder is, surprisingly, missing in the WAXS profile of PVDF/FM blend (Figure 3b, inset). This suggests that even if any polar phase is present in the blend, its mass fraction is small (i.e. below a few mass percent) and therefore it has no substantial influence on WAXS profiles shown in Figure 3. The lack of the shoulder at q~15 nm-1 in the WAXS profile of the PVDF/FM blend may additionally suggest that either the dimensions of polar crystals (ordered domains) present in the blend are in the range of few nanometers or their structure is highly defected.


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Figure 7. FTIR spectra recorded in the middle (MIR, a) and near (NIR, b) infrared region for the indicated samples with assignment of bands for the different , , and  phases of PVDF. For clarity, the spectra are vertically offset. The T-bars indicate the transmittance scale. All samples were crystallized at 150 °C. Peak assignments were adapted from Ref.54, with exception to the bands at 216 and 254 cm-1 that occur only in spectra of the samples where the polymorph was dominating phase.

Our WAXS and FTIR results (Figure 4 and 7 respectively) confirm that both OMT and OFM nucleate polar crystals in a similar manner as already reported by other researchers.55,

56

The

formation of one of the polar phases of PVDF was already evident from the WAXS analysis. The FTIR spectra of the PVDF/OMT and the PVDF/OFM nanocomposites shown in Figure 7 reveal characteristic bands corresponding to the - and - polymorphs of PVDF. According to the literature 54, the bands at 1274 cm-1 (Figure 7a) as well as 473 cm-1 and 443 cm-1 (Figure 7b) are related to vibrations characteristic of the -polymorph. Based on the same literature source54, we can attribute the bands at 1431 cm-1, 1230 cm-1, 810 cm-1 (Figure 7a) and also 481 cm-1 and 431 cm-1 (Figure 7b) to the - polymorph. The origin of the bands visible at 835 cm-1 and 510 cm-1 raises some controversies in the literature; however, they are related to formation to one of the polar polymorphs of PVDF: either the  or the . The presence of the bands characteristic of the polar polymorphs for PVDF and the absence of bands corresponding to the -polymorph indicate 27 ACS Paragon Plus Environment

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that both the PVDF/OMT and the PVDF/OFM nanocomposites contain predominantly the - or  polymorphs. A qualitative comparison of the FTIR spectra of PVDF/OMT and the PVDF/OFM nanocomposites leads to the conclusion that both polar phases, i.e., the - and -phase, are formed in the nanocomposites. However, transmittance intensities of the bands observed at 431 cm-1, 443 cm-1, and 473 cm-1 (Figure 7b) in addition to the bands at 810 cm-1 and 1230 cm-1 (Figure 7a) point towards a hypothesis that in the PVDF/OMT, the formation of  crystals is more preferred, while in the PVDF/OFM the -polymorph seems more dominating.


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Figure 8: (a-k) 19F MAS NMR spectra and (l-u) 19F MAS NMR spectra obtained with a long spin-lock pulse (20.0 ms) for the indicated samples (crystallization temperatures and PVDF to organoclay ratio by weight % recorded at 11.7 T and a MAS frequency of 25.0 kHz using high-power 1H decoupling during acquisition. Note how the large contribution from the amorphous phase, marked by an asterisk in (a-k), obscures most of the peaks associated with the crystalline phases for PVDF. The crosses indicate the position of PVDF defects associated with regio-irregular head-to-head and tail-to-tail structures.

To gain quantitative insights about the phase composition of the PVDF/OMT and PVDF/OFM nanocomposites, we have performed a series of solid-state

19F

magic-angle spinning (MAS)

NMR experiments with isotropic specimens as summarized in Figure 8. These experiments were obtained using direct 19F excitation (Figure 8a-k) and 19F excitation followed by a long spin-lock (SL) pulse of 20.0 ms, which selectively retains only magnetization from the crystalline domains (Figure 8l-u).57 Both experiments employed high-power 1H decoupling during acquisition. All observed 19F resonances fall in the chemical shift range between -80 and -120 ppm, while those from fluorine in the fluoromica are expected around -170 ppm to -180 ppm.58 Thus, we can assume that all observed 19F resonances in the range -80 to -120 ppm originate from the different phases of PVDF. Several resolved 19F resonances are observed in Figure 8l-u and a comparison with those obtained using direct 19F excitation show that these signals are largely hidden below the large and intense signal from the amorphous PVDF phase at -91.5 ppm. Previous work by Koseki et al. on PVDF and PVDF/PMMA films43 assigned the 19F chemical shifts for the three crystalline phases (-, -, and -) using

19F

spin-lock MAS NMR experiments and performed

their spectral assignment on the basis of gas-phase DFT shielding calculations for an 8-mer of PVDF as well as reference materials. From this work, the -phase was shown to include two 19F resonances (-79.4 and -93.7 ppm), the -phase four

19F

resonances (-79.4, -84.2, -93.7, and -

101.3 ppm), whereas the -phase only included a single 19F resonance at -93.7 ppm. We note that 29 ACS Paragon Plus Environment

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the 19F chemical shifts in Figure 8 are shifted by ~ -2 ppm compared to the work of Koseki et. al., which cannot be attributed to the Bloch-Siegert shift (~ -0.3 ppm).59 Most likely this difference is an effect of sample heating caused by the different MAS frequencies used here and in the work of Koseki et al.60 Similar results were found by Hucher et al., employing the SELDOM selective excitation technique to obtain the 19F MAS NMR spectra for the individual polymorphs.44 From a comparison of all 19F MAS NMR spectra obtained using the long spin-lock pulse (Figure 8l-u) at least six components are visible: two are readily attributed to the -phase from the pure PVDF spectra (-82.6 ppm and -96.1 ppm). The

19F

resonance at -96.1 ppm is attributed to the -phase

with some overlap with the -phase. This leaves four 19F signals for the -phase: -88.6, -92.3, 103.9 and a broad shoulder at ~ -100 ppm. Considering the previous assignment from Koseki et al.43 and Hucher et al., one would expect at least one 19F signal at the same shift as the - and phase.44 Thus, the origin of the broad signal at ~ -100 ppm remains ambiguous as this may contain the fourth resonance from the -phase. Alternatively, some form of polymer-clayinterface region could also be the source of these peaks. This leaves us with two possible spectral assignments (Model 1 and 2) as shown in detail in Figure 9, where each assignment results in slightly different amounts for the -phase. Furthermore, we assign the

19F

resonance at ~-

114 ppm to polymer defects corresponding to regio-irregular head-to-head and tail-to-tail structures.42


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Figure 9: 19F MAS NMR spectra recorded with a spin-lock time of 20.0 ms showing the phase composition based on (left) Model 1 and (right) Model 2. Table 1 summarizes the deconvolution results.

To quantify the amount of -, - and -phases, we have performed a spectral deconvolution based on both assignments (Model 1 and 2) for all

19F

MAS NMR spectra obtained using the

long spin-lock experiment, see Figure 9, as summarized in Table 1. This assumes that all 31 ACS Paragon Plus Environment

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crystalline PVDF polymorphs exhibit similar relaxation behavior during the spin-lock pulse (T1,).57 Since the - and -phases (and possibly also the -phase) have one

19F

resonance in

common (at 96.1 ppm), the amount of the -phase in Model 2 was determined by subtracting the appropriate amount for the -phase (as determined from the second peak belonging to the phase). For Model 1, the average of the three other -phase resonances were subtracted as well. A Gaussian peak and a G/L peak were used to model the broad shoulder around -100 ppm present in all samples, containing the - and -phase. For Model 2 only the latter G/L peak was attributed to the -phase; however, the strong overlap between the two peaks introduced some error to the overall fit. The pure PVDF sample contains almost exclusively crystals of the -phase in good agreement with the results from FTIR and WAXS. The crystal phases of the MT and FM containing samples are also dominated by the -polymorph; however, a small fraction of - and -phase is present, where the precise amounts are subject to crystallization temperature and composition. PVDF/MT crystallized at 151 °C shows the highest content of the - and -phases of 14/21% and 30%, respectively, while the PVDF/FM samples crystallized at 146 °C only contains ~10% - and below 4% of the -phase; the -phase remains the dominant polymorph in all cases, followed by the -phase. These results are in agreement with observations from WAXS, but contrast to the FTIR data, which suggests that a larger amount of -phase should be observed for PVDF/FM. Overall the addition of FM and MT leads to the formation of polar crystal phases with MT having a stronger effect. As has been observed by both X-ray diffraction and FTIR, the organoclay containing PVDF samples (PVDF/OFM and PVDF/OMT) crystallize predominantly as the polar crystal polymorphs (- and -). The low intensity of the peak at -82.6 ppm attributed to the -phase in comparison to the rest of the 19F NMR spectra in Figure 9 clearly confirms this behavior. The three peaks unambiguously assigned to the -phase are by far the most intensive 32 ACS Paragon Plus Environment

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for the PVDF/OFM and PVDF/OMT samples independent of the crystallization temperature with up to 64% of the -polymorph. Depending on the model chosen for deconvolution, the amount of -phase lies between 10% and 25%. Only minor amounts of -phase are detected.

Table 1: PVDF phase composition determined from spectral deconvolution of 19F MAS NMR spectra in Figure 9.

Phase compositiona) Samplea)

Crystallization temperature (°C)

-phase (%)

-phase (%)

Defects & unknown (%) M1 M2

-phase (%)

M1

M2

M1

M2

M1

M2

PVDF

144

90

90

7

7

0

0

3

3

PVDF

150

89

89

8

8

0

0

3

3

PVDF/MT

146

68

68

10

13

13

13

8

5

PVDF/MT

151

41

41

14

21

8

146

87

87

10

10

30 0b)

15

PVDF/FM

30 0b)

3

3

PVDF/FM

151

75

76

8

10

9

12

8

2

PVDF/OMT

158

7

7

10

26

64

59

20

9

PVDF/OMT

162

6

6

11

27

62

61

22

7

PVDF/OFM

158

5

5

14

29

59

58

22

8

PVDF/OFM

162

5

5

13

28

60

59

22

8

a) Model 1 (M1) is based on the data previously reported by Koseki et

al.43,

while Model 2 (M2) refers to the

modified assignment of the -phase described in the text (see also Figure 9). b) The -phase was omitted for this sample since the intensity of the only resolvable

19F

signal was below 1%,

which amounts to less than 4% of the overall intensity.

Concluding this section, it is worthwhile to briefly comment on the spectroscopic analysis of the phase composition for PVDF in the analyzed systems. Our considerations related to the phase composition were based on experimental results from FTIR and NMR spectroscopies. The use of FTIR has a long tradition in analysis of PVDF, since FTIR spectrometers are commonly available in most laboratories. However, as already pointed out by previous authors,6 the anisotropy of


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such samples may lead to erroneous conclusions regarding the composition of the samples. Thus, to quantitatively analyze the phase composition, it is necessary to obtain isotropic, powder samples suitable for FTIR analysis, which is not always feasible like in the case of our work. Qualitatively, the results from FTIR and solid-state NMR were quite similar as they are both able to detect the formation of polar PVDF crystals in the studied systems. However, the advantage of solid-state NMR is that this technique enables precise information about the phase composition of the samples without ambiguities related to sample texture.

Conclusions

b) a) We have demonstrated that a conventional melt extrusion of poly(vinylidene fluoride) (PVDF)

with 5 wt.% of organophilized fluoromica (organofluoromica) can result in intercalated nanocomposites, containing spatially oriented, polar crystals of PVDF ( and  polymorphs). Our studied systems contained either of the two organoclays: organomontmorillonite with the aspect ratio of particles ~12 or organofluoromica with particles aspect ratio of ~27. Particles of both the organoclays revealed a tendency to orient along the polymer melt flow and they both effectively nucleated the polar crystals. In the case of the nanocomposites with organofluoromica, however, the crystalline lamellar nanostructure of PVDF was clearly anisotropic: the amorphous/crystalline stacks formed a remarkable, chevron-like pattern in the sample volume. Time-resolved X-ray scattering experiments indicated that such a characteristic orientation of lamellar crystals resulted mainly from the solidification of the oriented polymer melt that was enhanced by the oriented organofluoromica particles. While the orientation of the PVDF lamellar crystals was found to be related to the aspect ratio of the organoclay particles, the formation of the polar polymorphs was mainly a result of the interactions of PVDF with the surface of particles. Pure PVDF crystallized mainly into the 34 ACS Paragon Plus Environment

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polymorph. Increasing the crystallization temperature resulted in formation of a small fraction (~8% of the total crystalline phase) for the -polymorph. Crystallization of PVDF in the presence of the unmodified clays also resulted in the formation of mainly the -polymorph. In this case, however, increasing the crystallization temperature caused a more pronounced increase in the content of polar crystals, especially the -polymorph. A major increase in the content of PVDF polar crystals was observed in the systems with the organoclays. Unlike in the pure PVDF and the blends of PVDF with unmodified clays, formation of the polar crystals in the PVDF/organoclay nanocomposites was found to be almost independent of the crystallization temperature. The content of the -phase in these nanocomposites was always smaller than 7 % of the total crystalline phase, while the -:polymorph ratio was close to 1:2 irrespectively of crystallization temperatures or the aspect ratio of organoclay particles. Based on these results, we believe that the application of high-aspect ratio organoclay particles and crystallization under flow can be further developed to obtain functional, electroactive PVDF-based materials, where such oriented structures of polar crystals are necessary.

Associated content The supporting information including additional WAXS patterns TEM micrographs, DSC data, and deconvolution of the NMR spectra are provided to support the main content. Acknowledgements This work was initially funded through Marie Curie Intra European Fellowship within 7th EC Framework Programme through a grant PIEF-GA-2009-253521. Additional financial support for continuation of the research was provided by National Science Centre Poland (Grant No. UMO35 ACS Paragon Plus Environment

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2017/25/B/ST5/02869. DC acknowledges the funding through statutory subsidy of Advanced Materials Engineering and Modeling Group at Faculty of Chemistry, Wroclaw University of Technology. AK thanks to Dr. Agnieszka Kiersnowska and Mr. Lampros Papoutsakis for help with preparation of the samples during the stay at FORTH, Heraklion. Synchrotron experiments have been carried out at ID15, ESRF. The expert advice and assistance of Thomas Buslaps is gratefully acknowledged. The Authors thank to Mr. Gunnar Glaßer from Max Planck Institute for Polymer Research for excellent TEM and SEM micrographs.

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Table of Contents graphics


Formation of Oriented Polar Crystals in Bulk Poly(vinylidene fluoride

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