Nanoparticle induced piezoelectric, super toughened, radiation resistant, multi-functional nanohybrids

Abstract

We have developed multifunctional nanohybrids of poly(vinylidene fluoride-co-chlorotrifluoroethylene) (CTFE) with a small percentage of surface modified inorganic layered silicate showing dramatic improvement in toughness, radiation resistant and piezoelectric properties vis-à-vis pristine polymer. Massive intercalation (d₀₀₁ 1.8 → 3.9 nm) of polymer inside the nanoclay galleries and unique crystallization behavior of the fluoropolymer on the surface of individual silicate layer has been reported. Toughness in the nanohybrid increases more than three orders of magnitude as compared to pure CTFE. High energy radiation (80 MeV Si⁺⁷) causes chain session, amorphization and creates olefinic bonds in the pure polymer while the nanohybrids are radiation resistant at a similar dose. Nanoclay induces the metastable piezoelectric β-phase in CTFE, suitable for sensor and actuator application. Molecular level changes after irradiation and controlled morphology for smart membrane have been confirmed by using spectroscopy, sol–gel technique, surface morphology studies and in situ residual gas analysis.

---

Introduction

Multifunctional nanostructured hybrids have a diverse array of applications such as structural composites, sensors, and electronic devices that have demonstrated good promise in addressing and offering solutions in many different applications. Nanohybrids (one of the components must have its dimension on the nanoscale) enhance physical and chemical performance with significant improvements in mechanical properties, corrosion resistance, thermal performances, processability and lower environmental impact. In most cases, advanced composites are susceptible to various forms of damage e.g. unexpected mechanical overload including impact and fatigue, radiation, heat and chemical attack. Although a large variety of polymer-layered silicate nanocomposites have been described, the compatibility of inorganic and organic matrices at the nanoscale, especially in the interface region, leading to morphology driven phase transformation, still represents a major challenge. To address the lingering nanostructure and achieve superior mechanical and radiation resistant properties, we developed polymer nanoclay hybrids to generate multifunctional nanohybrids. The nanohybrids introduced here can be used to create multifunctional materials and devices (e.g., sensors and actuators, space application, radiation resistant coating, morphology-directed structural transformation with stiffer and tougher polymeric materials by design), including a new advanced hybrid architecture. This material offers super toughening, radiation resistant and piezoelectric material, suitable for sensor and actuator applications. The superior properties of the hybrid materials occur only in presence of two-dimensional nanoparticles which alter the crystalline phase and morphology of matrix polymer suitable for different applications.

Multifunctional nanostructured hybrids are of considerable interest for electromechanical devices, optoelectronics and radiation resistant materials. The properties of poly(vinylidene fluoride-co-chlorotrifluoroethylene) (CTFE) are relatively better than any other existing homofluoropolymers, and include remarkable resistance towards flame, chemicals, solvents, heat and oxidation.¹Copolymer in presence of side groups generates disorder in the highly crystalline homopolymer.² The structure of pure CTFE copolymer is similar to that of poly(vinylidene fluoride) homopolymer and exists in several crystalline forms: α, β, γ and δ phases, depending on the geometric chain configurations such as trans (T) or gauche (G) linkages.^1,3 The metastable β-phase (all trans) is piezo- and pyroelectric, suitable for electroacoustic transducer and electromechanical applications in general, but this phase is difficult to obtain in bulk.^4–6Copolymer content in the main chain generates defects and has usually low crystallinity/high ductility, modifies the symmetry of the polymeric chain and modulates both intramolecular and intermolecular forces so that properties such as melting point, glass transition temperature, crystallinity, stability, elasticity, permeability and chemical reactivity will be varied in a wide spectrum in comparison to PVDF homopolymer. Further, the fluorinated copolymers produce disorder in the macromolecular chains. Addition of surface-treated two-dimensional nanoclay in PVDF, so-called nanohybrids, shows the growth of the β-form in bulk and has been extensively utilized in an attempt to enhance the mechanical, physical and thermal properties.⁷

Swift heavy ion (SHI) irradiation is the most general and powerful method to bring effective and systematic changes in polymer structure–property relationships. Interaction of an ion beam with polymeric materials causes chemical and physical changes including chain session,⁸ crosslinking,⁹ free radicals,¹⁰ formation of double bonds,¹¹ structural changes¹²etc. and it depends on linear energy transfer (LET), a measure of the energy transfer to the material as an ionizing particle travels through it. Usually, a high energy ion beam (∼100 MeV) destroys the pure polymer film as the energy is much higher than the C–C bond energy (∼80 Kcal mol^-1). Chain scission occurs after bombardment with sufficient energy to produce an isolated ‘spur’ with low LET, while high LET generates a couple of ‘spurs’ which further coalesce, leading to a crosslinked structure. Recently, we have reported the radiation resistant behavior (∼80 MeV SHI irradiation) of PVDF in presence of organically modified nanoclay,¹³ while pure PVDF films become brittle when exposed to similar ion beams. The swelling experiments and gel fraction studies of copolymer indicate that chain session is a major occurrence in pure copolymer while crosslinking is a core phenomenon in nanocomposites, presumably due to parallel chain configuration of the long chain molecules on the surface of nanoclay discs.¹⁴

In this work, we develop a novel nanohybrid of copolymer of vinylidene fluoride with a small percentage of layered silicates of nanometer dimension. Unique nanostructures and structures have been worked out by studying XRD and electron diffraction patterns, revealing piezoelectric phase in presence of nanoclay. Morphology-driven significant improvement in mechanical responses has been demonstrated with a plausible mechanism. The effect of swift heavy ions has been studied with possible molecular level changes with and without the presence of nanoclay. A model has been given for the unique crystallization behavior in presence of nanoclay.

Experimental

Materials

Commercial powdered SOLEF 31008, a copolymer of vinylidene fluoride and chlorotrifluoroethylene, used for the study, termed ‘CTFE’, was obtained from Ausimont, Italy (MFI:15 g/10 min, @ 230 °C, 5 kg). The organically modified clay was used as a nanofiller (Cloisite 30B; Southern Clay Inc. USA; having cation exchange capacity ∼90 meq/100 g). The modifier used in this nanoclay is methyl tallow bis-(hydroxy ethyl) ammonium ion exchanged montmorillonite and the modifier content is ∼40%. The gallery spacing of the organically modified clay is 1.8 nm. The lateral dimension of nanoclay is ∼250 nm.

Sample preparation

The solution casting technique was used to prepare the CTFE nanohybrids with 4 and 8 wt% of nanoclay. Dispersion/premix of nanoclay and CTFE, were made using DMF solvent. The CTFE–DMF premix was prepared at 50 °C after stirring up for to 2 h to dissolve the polymer. The nanoclay–DMF dispersion was sonicated at 50 °C for 60 min for a better mixing of nanoclay into the solvent. Nanoclay dispersion was mixed with polymer solution and was stirred at 50 °C for 6–8 h for proper homogeneous mixing. Thin films of nanohybrids were prepared after pouring the solution into a Petri dish, followed by it drying at 50 °C for 12 h. Finally, the nanohybrids were kept in a vacuum oven for 48 h at 70 °C to remove the trace amount of solvent. Henceforth, the nanohybrids will be designated as NC4 and NC8 for 4 and 8 wt% of nanoclay in polymer matrix, respectively. Both the solution cast nanohybrids and the pure polymer were melt-pressed into thin film of ∼30 μm thickness of size 1 × 1 cm² in a compression molding machine at 190 °C under 5 tons of pressure for irradiation of swift heavy ions.

Swift heavy ions irradiation

High energy ion beam irradiation was carried out on CTFE and its nanohybrids using 80 MeV Si⁷⁺ ion in a vacuum of ∼10⁶ torr using the 15 UD Pelletron accelerator at the Inter University Accelerator Centre (IUAC), New Delhi. The ion fluence was taken in the range of 1 × 10¹⁰ to 5 × 10¹² ions cm⁻² maintaining 0.5 pnA (particle nanoampere) beam current with a spot size of 2 × 2 mm². The electronic energy ions (S_e) and nuclear energy loss (S_n) were ∼2.3 KeV nm⁻¹ and ∼1.8 eV nm⁻¹, respectively, as estimated by SRIM calculation. Film thickness of the samples was chosen so that no ion will be implanted with the energy of SHI.

Residual gas analysis

Residual gas analysis was performed in situ during irradiation by using a quadrupole mass analyzer (model HAL 2/201, Hiden, UK).

Characterization

X-ray diffraction (XRD). The crystal structure was examined using a Bruker AAXS D8 Advance wide-angle X-ray diffractometer with Cu kα radiation with a graphite monochromator (wavelength, λ = 0.154 nm).

Differential scanning calorimeter (DSC). The thermal properties of the samples were measured using a Mettler 832 DSC instrument. The melting temperature, crystallization temperature and heat of fusion of pure CTFE, NCs and irradiated thin films were determined using a scan rate of 10 °C min⁻¹.

Morphological investigation. Transmission electron microscopy (TEM) was used to observe the nanoscale dispersion of nanoclay in the matrix polymer. TEM images were obtained using a Tecnai-G2-20 operated at an accelerating voltage of 200 kV. A thin layer was sectioned at −80 °C using a Leica ultracut UCT equipped with a diamond knife. The morphology of the pure CTFE and its nanohybrids (both irradiated and pristine) was investigated by using filed emission scanning electron microscopy (SEM) and atomic force microscopy (AFM). The surface morphology of the thin film by SEM was examined with a FESEM ZEISS SUPRA™ 40 instrument operated at 5 kV. All the samples were palladium coated by means of a sputtering apparatus before observation in SEM. An NT-MDT multimode atomic force microscope (Russia) controlled by a Solver scanning probe, was used in tapping mode with the tip mounted on a 100 μm long single beam cantilever with resonant frequency in the range of 240–255 kHz and corresponding spring constant of 11.5 N m⁻¹.

Mechanical properties. Tensile testing of injection molded samples was carried out by using an Instron 3369 tensile tester at a strain rate of 5 mm min⁻¹ at room temperature. Dynamic mechanical properties of the samples were studied with a dynamic mechanical analyzer Q 800 (TA instruments) in the tensile mode. The dynamic responses were tested from −80 °C to 150 °C at a frequency of 1 Hz with strain amplitude of 15 μm at the heating rate of 3 °C min⁻¹.

X-ray photoelectron spectroscopy . To identify the chemical bonding of pure polymer and its nanohybrids, an X-ray photoelectron spectroscopy (XPS) study was carried out at a base pressure greater than 5 × 10⁻⁹ Torr. Al-K_α radiations with incident energy 1486.6 eV were employed for the analyses. Graphitic C-1s peak (284.6 eV) has been used as an internal reference to correct the shifts in binding energies of core levels due to charging effect.

Sol–gel analysis. The sol–gel analyses of irradiated polymer and nanohybrids were conducted by immersing the films in DMF for 48 h at 60 °C followed by filtering and drying to determine the insoluble fraction.

Spectroscopy . UV-visible absorption spectra of the transparent thin films, before and after irradiation, have been taken by using Shimadzu (UV-17001) spectrophotometer operating in the spectral range of 200–1100 nm. Photoluminescence was measured using a Horiba Jobin Yvon photoluminescence spectrometer. FTIR spectra of the transparent thin film were made using a FTIR (JASCO-470 plus, Shimadzu-1400) in the range of 400–4000 cm⁻¹.

Results and discussion

Nanostructure and structure

The nanohybrids containing 4 and 8 wt% of organically modified nanoclay were prepared via a solution route. Fig. 1(a) shows the nanostructures of nanoclay and representative CTFE nanohybrid with 8 wt% nanoclay (NC8). The d₀₀₁-spacing of pristine organically modified nanoclay of 1.8 nm has significantly been expanded to 3.9 nm in the nanohybrid due to insertion of polymer inside the layered galleries. The extent of intercalation in CTFE is much larger compared to other fluoropolymers, e.g. pure PVDF (2.9 nm)¹³ and its copolymer with hexafluoropropylene (3.0 nm),¹⁴ using a similar quantity of nanoclay and the reason lies in the better interaction between the organically modified nanoclay and the more ionic chlorine substituted CTFE fluoropolymer. Interestingly, a strong peak at 2θ ∼ 6.1°, corresponding to a d-spacing of 1.45 nm, appears for nanohybrids irrespective of nanoclay concentrations. The possibility of it being the (002) plane is nullified both from the d-spacing point of view and higher intense than that of the (001) plane, instead reflection from the crystallized polymers on two opposite faces of the nanoclay layers has been described in the inset of the figure. It is worthwhile to mention that an individual inorganic layer has a thickness of ∼0.7 nm and polymer chain crystallized on both the surfaces of the silicate layers, hence validating the spacing of 1.45 nm. For pure PVDF and HFP copolymer, the peak positions were matched with d₀₀₂ spacing even though the intensity was higher. The nanostructure is unique as it is possible to differentiate the existing typical nanostructure in presence of nanoclay from the (002) plane while for PVDF (homopolymer) nanohybrid, the typical nanostructural peak in XRD exactly match with (002) plane which couldn't be demarcated as separate peak. It so happened that nanostructural peak is in a different position to that of the (002) plane arising from layered spacing of polymer intercalated nanoclay. The peak at 2θ ∼ 6.1° is also present in PVDF nanohybrids, whereas this peak is not observable in nanocomposites of other polymers like PLA, PCL or PP. In order to check the collapsing/degradation, we heated the nanoclay at 180 °C for 2 h and found that the (001) peak remains in the same position (ESI Fig. S1†). This work has unique evidence where the XRD peak has substantially deviated from the (002) position and has sufficient intensity arising from the greater coherency of the crystallized polymer on nanoclay surfaces to establish the said reflection precisely for the first time. Fig. 1(b) (image I) shows bright field (BF) transmission electron micrograph of the nanohybrid, showing good dispersion of nanoclay in the polymer matrix. Interestingly, electron diffraction patterns of the nanohybrid are very diverse in different locations indicating pure β-phase (110/200) plane in the vicinity of the nanoclay (image II), a mixture of (020) and (110) planes in α-phase and (110/200) β-phase near around nanoclay cluster (image III) and complete amorphous phase in between two nanoclay clusters, placed sufficiently apart from each other (image IV). Electron diffraction patterns have also been given in ESI Fig. S2† for better clarity. A schematic has been presented based on the structural aspect arising from the distribution of nanoclay and nucleation of particular crystalline form (image V). It is evident from the electron diffraction patterns that the piezoelectric β-phase nucleates on the surface of the nanoclay presumably due to close similarity of the structure, while some α-phase nucleates in the absence/loss of clay/solid heterogeneous support (as observed in image III) and the usual amorphous region mostly prevails in the portion far away from nanoclay stacks. However, the circumstance is like a crystalline regime (β-phase on the nanoclay vicinity and mixed α- and β-phase on top of that) grown on layered silicate and whole solid/crystalline phase makes an island structure covered with liquid amorphous zone (Fig 1(b), Scheme V). It is noteworthy to mention that predominantly β-phase has been crystallized on the surface of the nanoclay, suggestive of the XRD peak at ∼6.1° in Fig. 1(a), while a meager α-phase has also been crystallized on top of the β-phase, as observed in electron diffraction pattern. The structural change in presence of a small weight percentage of nanoclay has also been confirmed through FTIR, X-ray and DSC studies. The α-peaks at 490, 615, 763 and 976 cm⁻¹ in FTIR spectra^15,16 of pure CTFE disappear in nanohybrid and those of β-peaks¹⁵ at 510, 600 and 840 cm⁻¹ appear in presence of nanoclay, demonstrating the conversion of α-phase of pure CTFE into piezoelectric β-phase in nanohybrids (Fig. 1(c)). Another peak at 1275 cm⁻¹ for β-phase in NC4 has been presented in ESI Fig. S3† against no peak for pure CTFE in the same wavenumber region. This transformation of structure in presence of nanoclay has also been confirmed from XRD and DSC measurements given in the ESI (Fig. S4 (a–c)†) where the characteristic α-peaks for (100), (020), (110) and (120) crystalline planes of pure CTFE transform to (200)/(110) planes in nanohybrids corresponding to β-crystalline phase.¹³ Deconvoluted XRD pattern of NC4 shows 30% β-content in the nanohybrid. Further, the melting temperature of pure CTFE at 167 °C has increased to 171.5 °C for NC8, in addition to a significant decrease of heat of fusion of nanohybrid (∼15 J g⁻¹) as compared to pure CTFE (34 J g⁻¹), which also supports the structural change in presence of nanoclay.^13,14

Fig. 1 (a) XRD patterns of nanoclay and NC8, schemes embedded in the graph to show the origin of diffraction peaks; (b) (I) bright field transmission electron micrograph and electron diffraction pattern showing (II) β-phase, (III) mixed α- and β- phase, (IV) amorphous phase, corresponding to the specified selected regions as indicated and (V) a schematic representation of nanoclay dispersion and locations of various crystalline structures; (c) FTIR spectra of CTFE, NC4 and NC8 showing indicated α and β peaks.

Nature of interactions

The nature of interactions between the polymer and nanoclay is crucial to understanding the structural change of CTFE in the presence of nanoclay. We considered the extent of interaction through X-ray photoelectron spectroscopy (XPS) and UV-vis studies. All core level XPS spectra were compared to the C 1s neutral carbon (graphite) peak at 284.6 eV. Fig. 2(a) shows the C 1s spectrum for pristine CTFE and NC4 having three types of carbon after deconvolution: neutral graphitic carbon peak at 284.6 eV, exactly same for both CTFE and NC4; carbon bonded to hydrogen (CH₂) at 286.0 eV of pure CTFE has been shifted to 285.7 eV for NC4¹⁷ and carbon bonded to fluorine (CF₂) at 290.1 eV of pristine CTFE has shifted to 289.5 eV for NC4.¹⁸ Therefore, we observe a significant shift of peak position of 0.6 eV for CF₂ carbon and 0.3 eV for CH₂ carbon in NC4vis-à-vis pure CTFE towards lower energy due to significant interaction between dipole (C^δ+–F^δ−) of CTFE chain and nanoclay.¹⁹ The higher shift in CF₂ carbon as compared to CH₂ carbon is explained from the greater charge separation in the former case in attachment with more electronegative fluorine, which in turn is responsible for the whole interactions between nanoclay and matrix polymer through dipolar interactions. It is noteworthy to mention that the interaction reduces the partial positive charge on carbon atom in the nanohybrid (NC4), leading to the shifting of peak position to a lower energy range. XPS patterns of F 1s show the peak positions of 688.2 and 687.4 eV for pure CTFE and nanohybrid, respectively (Fig. 2(b)). Fluorine has more a partial negative charge due to a hydrogen bonded arrangement with the clay modifier, thereby justifying the shifting of peak position to lower energy range in nanohybrid.²⁰ Hence, XPS studies confirm the extensive interactions between nanoclay and polymer, which is a pre-requisite for structural change to arise. Further, UV-vis absorption spectra of nanohybrids with varying nanoclay concentrations have been presented in Fig. 2(c), showing a systematic blue shift with decreasing nanoclay content in nanohybrids. The optical absorption at 244 nm appears due to electronic excitation of olefinic group present in organic modifier (tallow) while no such peak appears for pure CTFE being linear fluoro-alkane (inset of Fig. 2c).

Fig. 2 Comparison of pure CTFE and its nanohybrids (a) XPS spectra of carbon 1s core level for CTFE and NC4 including deconvoluted peaks, (b) XPS spectra of fluorine 1s core level for CTFE and NC4, and (c) UV-vis absorption spectra of indicated nanohybrids. Inset figure shows the absorption spectra of pure CTFE and pristine organically modified nanoclay.

However, the peak position gradually shifted to lower wavelength (blue shift) with decreasing nanoclay concentration, presumably due to the subtle configurational changes of the organic modifier on top of inorganic nanoclay. The organic modifiers are squeezed upon crystallization of CTFE chains on the surfaces of nanoclays. As a result, the coated (covered with crystallized chain) and constricted organic modifiers need much a higher energy to be excited, resulting in blue shifting of the peak. The origin of blue shift usually arises when long chain molecules form aggregates,²¹ or complexes²² or are sterically hindered²³ in their chemical environment. On the contrary, red shift is reported with same nanoclay where there is no crystallization on the surface of the nanoclay.²⁴ Moreover, the intensity of the peak decreases with decreasing the concentration of nanoclay and is essentially due to low abundance. For comparison purposes, we have prepared with unmodified MMT clay which exhibits much less β-content as observed in FTIR studies, presumably because of less interaction (ESI Fig. S5(a)†). The interaction with 30B clay is considerably higher primarily due to two hydroxyethyl groups present in the organic part of the nanoclay. A schematic representation of α- and β-phase crystallization on MMT clay and organically modified 30B clay has been shown in ESI Fig. S5(b).† Nonetheless, the combined study of XPS and UV absorption demonstrate the cause of structural alteration through extensive interaction and constricted conformation of organic modifier attached to inorganic nanoclay.

Mechanical responses

The mechanical responses under uniaxial elongation of pure CTFE and NC4 have been presented in Fig. 3. The elongation at breaking point significantly increases from 65% for the pristine CTFE to 750% in NC4 nanohybrid, offering the hybrid toughness, as calculated from the area under the stress–strain curve, ∼1100% higher than that of the neat polymer. The Young's modulus of nanohybrid improves to 0.36 GPa from its value of 0.30 GPa in pristine CTFE. In addition, tensile strength of nanohybrid also improves to 33.4 MPa from its value of 29.6 MPa for pure CTFE. In conclusion, there are no trade-offs and there is an all round improvement of mechanical properties associated with the nanohybrid, which is usually not observed for most reported polymer-layered silicate nanohybrids. The super toughening phenomenon, more than three orders of magnitude, is presumably morphology-driven in presence of nanoclay. Compact and organized spherulites in pure CTFE completely disappear in the nanohybrid (Fig. 3, inset POM images) causing loss of birefringence in the mesh-like β-crystalline structure. This is further elucidated in inset SEM images where the mesh-like element is prominent in the nanohybrid instead of compact spherulite in pristine CTFE. The compact and highly crystalline spherulites are brittle, causing lower elongation at breaking point (65%) for pure CTFE. On the contrary, the tiny mesh-like structures in the nanohybrid appear from the crystallization of polymer on top of nanoclay as demonstrated in scheme V of Fig. 2b, help rotating and, thereby, dissipate energy from the crack tip front under stress field, building dramatic enhancement of toughness.^25,26 This is a unique report of improving such a significant improvement in toughness for a fluoropolymer in general against the literature reported 700% improvement in PVDF/layered silicate system.²⁶ Possible reasons include enhanced specific surface area and better mechanical interlocking/adhesion at the interface of nanofiller and the more polar chloro-substituted fluoropolymer, as evident from the greater toughness of graphene filled epoxy in comparison to CNT filled epoxy composite.²⁷ Much higher toughness of 2000 times is reported for PVDF by grafting with dimethylaminoethyl methacrylate and that improvement is not caused by structural change but by the chemical modification.²⁸ Further, the absence of a relaxation peak in dynamic measurement for nanohybrid at 100 °C, presumably due to chain movement inside crystalline zone²⁹ in contrast to pure CTFE, aggravated the better energy dissipating mechanism, as shown in Fig. 4 and ESI Table 1.† The NC4 exhibits higher storage as well as loss modulus. Pure CTFE shows three relaxations at −32 °C, 10 °C and 98 °C for glass transition, crystal–amorphous interface and within the crystalline region (pre-melting), respectively. The relaxation peak at 98 °C is absent in the nanohybrid indicating greater movement in nanohybrid due to tiny mesh-like β-crystallite causing remarkable toughness. The lower temperature transition around −32 °C occurs due to cooperative motions within the main chain in the amorphous phase, whereas the second transition is observed due to relaxation of crystal–amorphous interface. However, super toughened, stiffened and stronger nanohybrids are produced in presence of two-dimensional organically modified nanoclay.

Fig. 3 Stress-strain curves for pure CTFE and nanohybrid (NC4) showing the dramatic increase in elongation at break, embedded images (a₁), (a₂) show optical micrographs of CTFE and NC4 and (b₁), (b₂) show SEM images of CTFE and NC4, respectively.

Fig. 4 Dynamic mechanical response of pure CTFE and NC4 nanohybrid as a function of temperature. Storage modulus (E^′), loss modulus (E^′′) and loss tangent (tan δ) are plotted against temperature while maintaining the frequency 1 Hz constant. The heating rate was kept at 3 °C min⁻¹.

Radiation resistant behavior and molecular phenomena

Thin specimens (∼30 μm) were bombarded with swift heavy ions (SHI; 80 MeV Si⁺⁷) at the Inter University Accelerator Centre, New Delhi, using a 15UD Pelletron accelerator operated at ∼10⁶ torr to perceive the effect of high energy radiation on polymers. The ion fluence range was taken between 1 × 10¹⁰ to 5 × 10¹² ions cm⁻² and 0.5 pnA (particle nanoampere) beam current was maintained with a spot size of 2 × 2 mm². The electronic energy ions (S_e) and nuclear energy loss (S_n) were ∼2.3 KeV nm⁻¹ and ∼1.8 eV nm⁻¹, respectively, as estimated by SRIM calculation. Thin film thickness of specimens and energy of Si⁷⁺ were so chosen to ensure that no ion will be implanted. Surface morphology through 3D- AFM images indicate the significant roughening of surface for pure CTFE under SHI irradiation, while it is remarkably suppressed in nanohybrids under the same irradiation conditions (Fig. 5(a)). The height profiles (uniform roughness) of pure CTFE and NC4 before and after SHI irradiation quantitatively measure the surface roughness at different fluences. Gradual increase of surface roughness is noticed for pure CTFE with increasing fluence while more or less similar surface roughness is observed for NC4, showing prohibition of surface damage in nanohybrids in the presence of nanoclay. The heat of fusion (ΔH) of pure CTFE has drastically been reduced after SHI irradiation and that reduction further decreases at higher fluences (Fig. 5(b)), presumably due to chain session and amorphization under SHI irradiation.³⁰ On the other hand, there is minimal reduction in ΔH for the nanohybrid under the same conditions of SHI irradiation, strongly suggesting the non-destructive nature of polymer matrix in presence of nanoclay. It is presumed that disc-like inorganic nanoclay protects the matrix polymer from high energy radiation. Further, the gel fraction, a measure of crosslinking, increases with increasing fluence both for pure CTFE and nanohybrids but the relative rate is considerably higher for the nanohybrid in comparison to pure CTFE (Fig. 6). Gel fraction increases up to 94% for nanohybrid as compared to 80% for pure CTFE at 5 × 10¹² ions cm⁻² fluence. In general, SHI irradiation originate chain scission, crosslinking and permanent free radicals in polymer chains.³¹ In the presence of modified nanoclay, crosslinking phenomenon is favored with all trans β-crystalline polymer chains due to recombination of radicals (formed during SHI bombardment) on the proximity of 2-D solid layered silicates while chain session is a prominent phenomenon in pure CTFE. A plot of vs., according to classical Charlesby-Pinner equation,³² where s = 1 − F is the soluble fraction, and D is the dose, shows greater solubility fraction for pure CTFE as compared to the nanohybrid. The ratios of , calculated from the extrapolation of the linear part of the curves at = 0 (D → ∞), are 1.75 and 1.0 for pure CTFE and NC4, respectively, where G_s and G_c are the radiation–chemical yields of chain session and crosslinking, respectively. The higher value for pure CTFE clearly indicates the greater chain session and less crosslinking while the polymer chains in nanohybrids are significantly crosslinked due to an all trans β-chain conformation on solid template of nanoclay. The swelling experiments and gel fraction studies strongly indicate that chain session is a major occurrence in pure CTFE while crosslinking is a core phenomenon in nanohybrids presumably due to parallel chain conformation of CTFE molecules on the surface of nanoclay discs. Hence, the nanohybrids of CTFE are radiation resistant and we outlined a plausible reason of the remarkable behavior.

Fig. 5 (a) AFM topographs of CTFE and NC4 thin film before and after SHI irradiation at indicated fluence with height profile. (b) Heat of fusion as a function of fluence for pure CTFE and NC4. Short horizontal solid and dashed lines are the values for the corresponding pristine non-irradiated samples indicated as ‘0’.

Fig. 6 Gel fraction of pure CTFE and NC4 nanohybrid in the swelling experiment as a function of fluence. The solubility parameter vs. fluence graph is embedded in the inset.

The changes in the molecular level as a consequence of SHI irradiation have been shown in Fig. 7. An absorption peak at 223 nm appears for pure CTFE after irradiation which gets more intense with increasing fluence presumably due to evolution of olefinic bond as a result of the cleavage of the C–Cl bond and subsequent hydrogen radical, leading to release of HCl gas. The stronger peak arises with higher fluence as the release of HCl is prominent, leading to formation of more double bonds. Nanohybrid NC4 shows an absorption peak at 236 nm even before irradiation, owing to double bond present in the organic modifier, and that peak becomes more intense with increasing fluence (Fig. 7(a)). The absorption peak for separate double bonds in the nanohybrid is not observed and it is presumed that the peak may merge with the modifier band causing a slight intense band. Using the sol–gel technique, we established that crosslinking is prominent in nanohybrid (greater than pure CTFE) and UV-vis clearly suggests that generation of double bonds is prominent in pure CTFE, while it is negligibly small or not appears in nanohybrids. The evidence of emission from those double bonds, created by SHI irradiation in pure CTFE, is apparent from the moderately strong peaks at 338 nm in photo luminance studies (Fig. 7(b)). A very weak peak for the nanohybrid at the same peak position clearly suggests the minimal abundance of double bond/emission behavior vis-à-vis pristine CTFE. It has to be mentioned that both pure CTFE and nanohybrid do not exhibit any emission peak in that region before SHI irradiation. However, the double bond has been created after releasing HCl from the polymer chain under SHI irradiation. The surface morphology of pure CTFE exhibits lot of pores after irradiation, while the number of pores is appreciably controlled in nanohybrids (Fig. 8(a)). Moreover, the relative size of the pores is comparatively less in the nanohybrid as compared to pristine CTFE (ESI Fig. S6†). During irradiation, the polymer is in a melted state, as SHI raises temperature of the substrate due to linear energy transfer, and the produced HCl excreted from the film surface causes holes in the film. Pure CTFE creates large amounts of HCl upon irradiation as the double bond formation is predominant, while a lesser amount of HCl is formed in nanohybrids as the double bond formation is remarkably suppressed in the crosslinking dominated phenomenon in the presence of nanoclay. It is worth noting that fibrillar morphology in pure CTFE is a part of spherulites which is missing in nanohybrids, as discussed earlier. The pores in the polymer film help them to impermeate gases through the film and that can be controlled/restricted in nanohybrids due to the regulated size and number of the pores. Further, with the regulated pore size, the material has the potential to act as a smart membrane as the matrix is piezoelectric, as revealed earlier. Moreover, the evolution of HCl gas has been confirmed by measuring residual gas analysis in situ during SHI irradiation. Residual gas analysis was performed in situ using quadrupole mass analyzer (model HAL 2/201 from Hiden, UK). The evolution of HF and HCl gas as a function of time has been shown in Fig. 8(b). Interestingly, the generation of gases are considerably suppressed in nanohybrids vis-à-vis pure CTFE and, thereby, further strengthen the spectroscopic and morphological evidence caused by swift heavy ions.

Fig. 7 (a) UV-vis spectra of pure CTFE and NC4, and (b) photoluminance spectra of CTFE and NC4 before and after irradiation at indicated fluences. Specimens before irradiation are marked as ‘0’.

Fig. 8 (a) SEM images of CTFE (left) and NC4 (right) at the fluence 1 × 10¹⁰ ions cm⁻², (b) variation of evolution of HF and HCl gas for CTFE and NC4 as a function of time (in situ) during swift heavy ion irradiation using quadrupole mass analyzer.

The utility of this nanohybrid is demonstrated across various platforms, which shows the benefits of working with multifunctional materials in composite systems. The developed nanohybrids can be used as electromechanical devices and smart membranes, which can be operated in any adverse environment being radiation resistant and super toughened material with a unique nanostructure. This work hopes to serve as a guide to researching the behaviors of multifunctional materials and stimulating further research on fluoropolymer composites and related device fields.

Conclusion

Piezoelectric, radiation resistant and super toughened nanohybrids have been prepared by dispersing a small weight percentage of organically modified nanoclay in fluoropolymer. Unique nanostructure, crystallization of thin polymer on the surface of the nanoclay, has been revealed through XRD and TEM. The reason for typical crystallization has been explored and schematic has been shown with evidence from morphology and electron diffraction studies explaining super toughening phenomena in nanohybrids. Swift heavy ions generate pores and that can be controlled (shape and size) in nanohybrids, leading to the opportunity to prepare smart membranes, as the matrix is a piezoelectric phase as evident from structural characterization. Molecular level changes (e.g. formation of double bond and crosslinking) have been explored by means of in situ gas evolution, spectroscopic and sol–gel techniques before and after SHI irradiation.

Acknowledgements

The authors acknowledge the receipt of research funding and peletron source from Inter University Accelerator Centre (IUAC), New Delhi (Project No. IUAC/XIII.7/UFUP-3932/5603). The authors also acknowledge Dr D. K. Avasthi, Saif A. Khan and Pawan K. Kulriya for assisting in shift heavy ion measurement and XRD studies. The kind supply of HFP samples by Ausimont, Italy is highly acknowledged. Vimal K. Tiwari acknowledges the award of Senior Research Fellowship of CSIR, India.

References

1. Bruno Ameduri, Chem. Rev., 2009, 109(12), 6632–6686.
2. Baojin Chu, Xin Zhou, Kailiang Ren, Bret Neese, Minren Lin, Qing Wang, F. Bauer and Q. M. Zhang, Science, 2006, 313, 334–336.
3. Zhimin Li, Yuhong Wang and Z.-Y. Cheng, Appl. Phys. Lett., 2006, 88, 062904-1–062904-3.
4. J. Wang, H. Li, J. Liu, Y. Duan, S. Jiang and S. Yan, J. Am. Chem. Soc., 2003, 125, 1496–1497.
5. J. Scheinbeim, C. Nakafuku, B. A. Newman and K. D. Pae, J. Appl. Phys., 1979, 50, 4399–4405; B. A. Newman, C. H. Yoon, K.d. Pae and J. Scheinbeim, J. Appl. Phys., 1978, 49, 4601–4603.
6. A. J. Lovinger, Polymer, 1981, 22, 412–413.
7. L. Priya and J. P. Jog, J. Polym. Sci., Part B: Polym. Phys., 2002, 40, 1682–1689; L. Priya and J. P. Jog, J. Polym. Sci., Part B: Polym. Phys., 2003, 41, 31–38.
8. E. Balazat, N. Betz and S. Bouffard, Nucl. Instrum. Methods Phys. Res., Sect. B, 1995, 105, 46–54.
9. B. Daudin, J. F. Legrand and F. Macchi, J. Appl. Phys., 1991, 70(8), 4037–4044.
10. E. Adem, G. Burillo, E. Munoz, J. Rickards, L. Cota and M. Avalos-Borja, Polym. Degrad. Stab., 2003, 81, 75–79.
11. N. Betz, A. Le Moël, E. Balanzat, J. M. Ramillon, J. Lamotte, J. P. Gallas and G. Jaskierowicz, J. Polym. Sci., Part B: Polym. Phys., 1994, 32, 1493–1503.
12. V. K. Tiwari, D. K. Avasthi and P. Maiti, ACS Appl. Mater. Interfaces, 2011, 3(5), 1398–1401.
13. V. K. Tiwari, P. K. Kulriya, D. K. Avasthi and P. Maiti, ACS Appl. Mater. Interfaces, 2009, 1(2), 311–318.
14. V. K. Tiwari, P. K. Kulriya, D. K. Avasthi and P. Maiti, J. Phys. Chem. B, 2009, 113, 11632–11641.
15. S. Lanceros-Mendez, J. F. Mano, A. M. Costa and V. H. Schmidt, J. Macromol. Sci., Part B: Phys., 2001, 3&4, 517–527.
16. A. Salimi and A. A. Yousefi, Polym. Test., 2003, 22, 699–704.
17. E. Adem, F. F. Castillon, G. Burillo, J. Rickards, E. Munoz and L. Cota, Polym. Bull., 2004, 52, 163–170.
18. N. Moussaif, C. Pagnoulle, J. Riga and R. Jerome, Polymer, 2000, 41, 3391–3394.
19. Paul Podsiadlo, Amit K. Kaushik, Elten M. Arruda, Anthony M. Waas, Bong Sup Shim, Jiadi Xu, Himabindu Nandivada, Benjamin G. Pumplin, Joerg Lahanm, Ayyalusamy Ramamoorthy and Nicholas A. Kotov, Science, 2007, 318, 80–83.
20. Benjaram M. Reddy, Biswajit Chowdhury and Panagiotis G. Smirnitis, Appl. Catal., A, 2001, 211, 19–30.
21. B. Vishalakshi, J. Polym. Sci., Part A: Polym. Chem., 1995, 33, 365.
22. H. L. Wang, M. Grigorova, E. S. Maniloff, D. W. McBranch and B. R. Mattes, Synth. Met., 1997, 84, 253–254.
23. Carine Edder, Paul B. Armstrong, Kris B. Prado and M. J. Frechet, Chem. Commun., 2006, 1965–1967.
24. A. Mishra, Biswa P. D. Purkayastha, Jagat K. Ray, Vinod K. Aswal and P. Maiti, Macromolecules, 2010, 43(23), 9928–9936.
25. D. Shah, P. Maiti, E. Gunn, D. F. Schmidt, D. D. Jiang, C. A. Batt and E. P. Giannelis, Adv. Mater., 2004, 16(14), 1173–1177.
26. D. Shah, P. Maiti, D. D. Jiang, C. A. Batt and E. P. Giannelis, Adv. Mater., 2005, 17(5), 525–528.
27. A. Mohammad Rafiee, Javed Rafiee, Zhou Wang, Huaihe Song, Zhong-Zhen Yu and Nihkil Koratkar, ACS Nano, 2009, 3(12), 3884–3890.
28. Sanjoy Samanta, Dhruba P. Chatterjee, Swarup Manna, Amit Mandal, Ashesh Garai and Arun K. Nandi, Macromolecules, 2009, 42, 3112–3120.
29. K. Nakagawa and Y. Ishida, J. Polym. Sci., Part A-2, 1973, 11, 1503–1533.
30. E. Adem, J. Rickards, G. Burillo and M. Avalos-Borja, Radiat. Phys. Chem., 1999, 54, 637–641.
31. E. H. Lee, Nucl. Instrum. Methods Phys. Res., Sect. B, 1999, 151, 29–41.
32. A. Charlesby and S. H. Pinner, Proc. R. Soc. London, Ser. A, 1959, 249, 367–389.

---

Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c1nr11009h

---