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Poly(Vinylidene fluoride-co-hexafluoro propylene)/Layered Silicate Nanocomposites: The Effect of Swift Heavy Ion Vimal K. Tiwari,† Pawan K. Kulriya,‡ Devesh K. Avasthi,‡ and Pralay Maiti*,† School of Materials Science and Technology, Institute of Technology, Banaras Hindu UniVersity, Varanasi 221 005, India, Inter UniVersity Accelerator Centre, Aruna Asaf Ali Marg, New Delhi 110 067, India ReceiVed: May 26, 2009; ReVised Manuscript ReceiVed: July 3, 2009
Poly(vinylidene fluoride-co-hexafluoropropylene) (HFP) nanocomposites with layered silicate have been synthesized via the melt extrusion route. The intriguing nanostructure, crystalline structure, morphology, and thermal and mechanical properties of the nanocomposites have been studied and compared critically with pristine polymer. HFP forms intercalated or partially exfoliated nanostructure (or both) in the presence of nanoclay, depending on its concentration. The bombardment of high-energy swift, heavy ions (SHI) on HFP and its nanocomposites has been explored in a wide range of fluence. The nanoclay induces the piezoelectric β-phase in bulk HFP, and the structure remains intact upon SHI irradiation. SHI irradiation degrades pure polymer, but the degradation is suppressed radically in nanocomposites. The heat of fusion of pristine HFP has drastically been reduced upon SHI irradiation, whereas there are relatively minute changes in nanocomposites. The coarsening on the surface and bulk of HFP and its nanocomposite films upon SHI irradiation has been measured quantitatively by using atomic force microscopy. The degradation has been considerably suppressed in nanocomposites through cross-linking of polymer chains, providing a suitable high-energy, radiation-resistant polymeric material. A mechanism for this behavior originating from the swelling test and gel fraction (chemical cross-linking) as a result of SHI irradiation has been illustrated. Introduction Swift heavy ion (SHI) irradiation is a special technique for inducing physical and chemical modification in bulk materials. It is notable to see modifications such as chain scission, crosslinking,1 and structural changes2,3 after SHI irradiation on polymers. Controlled SHI irradiation with fixed linear energy transfer (LET) is responsible for structural modification, latent track dimensions, and development of a membrane after etching. The chemical and physical changes (e.g., cross-linking, chain scission, electric and dielectric properties) of different polymers have been reviewed depending on LET.4 Chain scission occurs after bombardment with sufficient energy to produce an isolated “spur” with low LET, whereas high LET generates a pair of “spurs” which further coalesce, leading to a cross-linked structure. SHI has adequate LET that is responsible for higher energy modification than that induced by electron- or γ-beam and light ions having low LET. There are two fundamental processes of energy absorption during SHI irradiation: (1) nuclear stopping generated by the interaction with target nuclei and (2) electronic stopping as a result of interaction with target electrons, which dominates if the velocity of the irradiated ions is comparable to the Bohr velocity. Fluoropolymers such as poly(vinylidene fluoride) (PVDF) and its copolymers with hexafluoropropylene (P(VDF-HFP)) are highly nonreactive thermoplastics and are technologically important because of their availability in different crystalline forms. Fluoropolymers are semicrystalline and have applications in such diverse fields as paint for skyscrapers, transducers for sensitive scientific instruments, pipes for caustic chemical * To whom correspondence should be addressed. E-mail: pmaiti.mst@ itbhu.ac.in. † Banaras Hindu University. ‡ Inter University Accelerator Centre.
Figure 1. Wide angle X-ray diffraction patterns of organically modified clay, pure HFP, and its nanocomposites. The numbers in nanometers indicate the interplanar distance.
byproducts, and anticorrosive materials. The crystalline structure of the P(VDF-HFP) copolymer is similar to that of PVDF.5,6 The degree of crystallinity of the P(VDF-HFP) is sufficiently reduced in comparison to pure PVDF, whereas the flexibility and chemical resistance are enormously enhanced.6,7 PVDF exists in several crystalline forms: R, β, γ, and δ phases, depending on the geometric chain configurations, such as trans j ) is inactive (T) or gauche (G) linkages.8 The R-phase (TGTG with respect to piezo- and pyroelectric properties, but the β-form (all trans) exhibits ferroelectric activity, suitable for electroacoustic transducer applications.9 The β-form can usually be obtained from melt crystallization at high pressure,10 polling at high voltage,11 recrystallization of carbon-coated highly oriented ultrathin film,12 and molecular epitaxy on the surface of potassium bromide.13 Addition of nanoclay in PVDF, so-called nanocomposites, shows the growth of the β-form and has been extensively utilized in an attempt to enhance the mechanical,
10.1021/jp904907y CCC: $40.75 2009 American Chemical Society Published on Web 08/04/2009
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Figure 4. FTIR spectra of pure HFP and its indicated nanocomposites. Different peaks have been assigned for the R- and β-phases of HFP.
Figure 2. Bright field transmission electron micrograph of HFP nanocomposite (NC4). The individual layer and tactoids are indicated.
silicate. The nanostructures, crystalline structure, morphology, and thermal and mechanical properties of the nanocomposites have been demonstrated in comparison to pure polymer, raising the technological importance of the nanocomposites. The effect of swift, heavy ions on HFP and its nanocomposites has been evaluated in terms of radiation resistance, structure, morphology, and gel fraction. A mechanism of chemical changes that occurs in HFP (e.g., chain scission, cross-linking etc.) upon SHI irradiation has been elucidated. Experimental Section
Figure 3. Wide-angle X-ray diffraction patterns of pure HFP and its indicated nanocomposites before irradiation.
physical and thermal properties.14-17 This is the most convenient way to produce piezoelectric polymeric materials in bulk, reported for the first time by Priya et al.14 PVDF films become brittle when exposed to high-energy ion beams. Recently, we have reported the radiation-resistant behavior of PVDF in the presence of organically modified nanoclay after high -energy (∼80 MeV) SHI irradiation.18 There is an increase in crystallinity of PVDF after γ- and electron beam irradiation.19,20 In contrast, the degree of crystallinity reduces in PVDF upon high-energy ion irradiation.21 The possibility of the formation of latent ion tracks and graft copolymer with different monomers is made possible after bombarding polymer with suitable SHI. The inhomogeneous and localized modifications of polymer lead to new directed properties and are of interest in industrial applications concerning the preparation of biomembranes and sensor technology.22-25 The gas permeability for hydrogen and carbon dioxide of these membranes has been reported as a function of etching time.26 Permeability of membranes irradiated by using high fluence was higher than the membrane irradiated with low fluence. The electrical properties have been affected by the presence of residual ions, permanent dipoles, space charges, nonreactive gases, and oxidation byproducts.27 The effect of nanoclay on ion tracks can be of interest to regulate the size of pores. Moreover, different sizes of SHI can be used to generate different track dimensions, depending on the application. In the present study, we report the nanocomposites of poly(vinylidene fluoride-co-hexafluoro propylene) with layered
Materials. A commercial SOLEF 11008 copolymer of vinylidene fluoride and hexafluoropropylene used for the study was supplied by Ausimont, Italy. Henceforth, the copolymer will be termed as HFP. An organically modified clay, Cloisite 30B (bis-(hydroxyethyl)methyl tallow ammonium ion-exchanged montmorillonite), purchased from Southern Clay Inc., USA, was used as the nanofiller. Tallow is a mixture of C18 and C16 long-chain alkene. After the organic modification, the interplanar distance of the clay increases to 1.8 nm from the 1.1 nm spacing of unmodified sodium clay. The lateral dimension of nanoclay is ∼250 nm. Preparation of Nanocomposites. Initially, the copolymer beads were converted into powder form by using a home-built chip sizer. A requisite amount of HFP powder and nanoclay (4 and 8 wt %) were mixed well in a high-speed (1000 rpm) mixer before being put it into the extruder. The nanocomposites were prepared by the melt extrusion method. Extrusion was carried out in a twin-screw extruder (Hakke Mini Lab). The mixing was done at a temperature of 200 °C for 10 min under a shear rate of 100 rpm. Adsorbed water was removed by keeping the nanocomposites in a vacuum oven for 24 h at 60 °C. Henceforth, the nanocomposites will be designated as NC4 and NC8 for 4 and 8 wt % of nanoclay in polymer matrix, respectively. Both the extruded strips (nanocomposites) and 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. HFP and its nanocomposites were irradiated using 80 MeV Si7+ ion in a vacuum of ∼10-6 torr at the Inter University Accelerator Centre, New Delhi, India. To avoid heating effects, the beam current was maintained at 0.5 pnA (particles nanoampere) with a spot size of 2 × 2 mm2. The SHI beam was scanned by electromagnetic scanner in an area of 1 × 1 cm2. SHI irradiation was done with the representative HFP nanocomposite (NC4). The ion fluence range was kept between 1010 and 5 × 1012 ions/cm2. The electronic energy loss (Se) was ∼2.3 KeV/nm, whereas the nuclear energy
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Figure 7. Dynamic mechanical response of pure HFP and its indicated nanocomposites as a function of temperature in tensile mode: (a) storage modulus, (b) loss modulus, and (c) tan δ.
TABLE 1: The Storage Modulus of Pure HFP and its Nanocomposites at different Temperature Rangesa sample
Figure 5. DSC thermograms of pure HFP and its indicated nanocomposites: (a) melting endotherm, and (b) cooling curves before irradiation at a heating/cooling rate of 10 °C/min.
HFP NC4 NC8 HFP NC4 NC8 HFP NC4 NC8
temp (°C) -70
25 140
storage modulus/GPa 6.254 8.129 8.549 1.151 1.301 1.547 0.0787 0.1068 0.1374
% increase 30 36.7 13 34.4 36 75
a The percentage increment was calculated in comparison to pure HFP.
Figure 6. Polarizing optical images of pure HFP and its representative nanocomposite. Birefringence has not been observed for the nanocomposites, but distinct spherulites were noticed for pure HFP.
loss was ∼1.8 eV/nm, as estimated by SRIM simulation.28 Both the HFP and the NC’s were irradiated with 80 MeV Si+7 ions at various fluences. Thickness of the thin film and energy of Si+7 ions were chosen in such a way that the range (distance covered by the particles with that initial kinetic energy) of the ions is greater than the film thickness to ensure that no ion will be implanted rather will pass through the thin film. X-ray Diffraction (XRD). X-ray diffraction experiments were performed using a Bruker AXS D8 Advance wide-angle X-ray diffractometer with Cu KR radiation and a graphite monochromator (wavelength, λ ) 0.154 nm). The generator was operated at 40 kV and 40 mA. Thin film of pure HFP and NC samples (before and after ion irradiation) were placed on a quartz sample holder at room temperature and were scanned at diffraction angle 2θ from 1° to 40° at the scanning rate of 1°/ min to explore the nanostructure and crystalline structure and, thereby, the effect of nanoclay on structure of the matrix polymer. Differential Scanning Calorimeter (DSC). The melting temperature, crystallization temperature, and heat of fusion of
pure HFP, NC’s, and irradiated thin films were measured in a Mettler 832 DSC instrument. The samples were heated at a scan rate of 10 °C/min. The peak temperature and enthalpy of fusion were measured from the endotherm using a computer attached to the instrument. After the first melting, the samples were cooled down at a constant rate of 10 °C/min to find the crystallization temperature and heat of crystallization in a similar fashion. Further, a second heating was taken to ensure the amount of crystallinity and melting temperature after removing all the thermal history in the first run. The DSC was calibrated with indium and zinc before use. 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 Philips CM-10 operated at an accelerating voltage of 100 kV. A thin layer was sectioned at -80 °C using a Leica ultracut UCT equipped with a diamond knife. The morphology of the thin film of pure HFP and nanocomposites (both irradiated and pristine) was investigated by using scanning electron microscopy (SEM) and atomic force microscopy (AFM). The surface morphology of the thin film was examined with a LEO 435 VP instrument operated at 15 kV. All the samples were gold-coated by means of a sputtering apparatus before observation in SEM. A NT-MDT multimode atomic force microscope, Russia, controlled by a Solver scanning probe, was used for the surface morphology study. Tapping mode was used with the tip mounted on a 100-µm-long, single-beam cantilever with resonant fre-
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Figure 10. Heat of fusion as a function of fluence for pure HFP and indicated nanocomposite. The solid and dashed lines are a guide to the eye. The short horizontal solid and dash lines are the values for the corresponding pristine samples indicated as ‘0’. Figure 8. Wide-angle X-ray diffraction patterns of (a) pure HFP, and (b) NC4 at indicated fluences. The number “0” indicates the zero fluence or pristine polymer/nanocomposite (before irradiation).
were then cut into 65 × 5 mm2 dimension with the help of a die and punch arrangement. Determination of Gel Fraction and Swelling. The soluble component of irradiated polymer and nanocomposites was extracted by immersing the samples in DMF at 60 °C for 48 h, followed by filtering and drying to determine the insoluble part. The insoluble (or gel) fraction, F, was calculated from
F)
Figure 9. DSC thermograms of (a) pure HFP, and (b) NC4 at indicated fluences during first heating at a heating rate of 10 °C/min.
quency in the range of 240-255 kHz and corresponding spring constant of 11.5 N/m. A thin film of ∼30 µm was used for AFM studies for both the HFP and NC’s before and after irradiation. The average individual pitting diameter and depth was measured from around 30-40 separate pittings in the AFM micrographs. The morphology of the thin film (∼30 µm) in an optical range was examined using a polarizing optical microscope (POM) (Nikon) after crystallizing the samples at 140 °C up to full solidification on a Mettler hot stage. Dynamic Mechanical Analysis (DMA). Dynamic mechanical properties of the samples were studied on thick films of pure HFP and nanocomposites with a dynamic mechanical analyzer Q 800 (TA Instruments) in the tensile mode. The dynamic responses were tested from -80 to 150 °C at a frequency of 1 Hz with strain amplitude of 15 µm and at a heating rate of 3 °C/min. Samples having a rectangular cross section were used for DMA testing. HFP and NC’s were first compression-molded into films of ∼1 mm thickness. These films
mf × 100 mi
(1)
where mi and mf are the initial and final weight, respectively. The swelling experiment was performed by measuring the weight of the polymer dipped into DMF at 60 °C at different time intervals. It is worth to mentioning that HFP is readily soluble in hot DMF, but SHI irradiated HFP and its nanocomposites exhibit some remnant, even after 48 h of exposure in DMF, due to cross-linked structure. Molecular Weight Determination. The molecular weights of pure HFP and NC’s were measured by using GPC (YoungLin Instrument Co.) in DMF as eluent at 70 °C using polystyrene standard. Both irradiated and pristine samples were dissolved in DMF for 48 h at 60 °C, and the extracted solution was taken for the molecular weight measurement. Results and Discussion Nanostructure. The X-ray diffraction patterns of organically modified nanoclay, pure HFP, and the nanocomposites with 4 and 8 wt % clay contents are shown in Figure 1. The (001) diffraction peak of nanoclay represents the interplanar distance of silicate layers of 1.8 nm (2θ ∼ 4.8°). The disappearance of the (001) diffraction peak in NC4 might be due to the disordered structure of the layered silicates, and the same peak appears at 2.9°, corresponding to the d001 ∼ 3 nm for NC8, suggesting an ordered intercalated structure with the insertion of a polymer chain in the nanoclay galleries.29 There is another intense diffraction peak at 2θ ∼ 6.15° for both NC4 and NC8 that corresponds to a d-spacing of 1.45 nm. The peak position exactly matches with higher order (002) planes, whereas its high intensity might be due to the additional factors discussed in the structural section. Figure 2 shows the bright field TEM image of melt-extruded NC4 sample in which the clay platelets are nicely dispersed in the HFP matrix, exhibiting discrete layers of nanoclay as well as some tactoids. The TEM image supports
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Figure 11. AFM images of HFP and NC4 thin film before and after SHI irradiation at indicated fluence obtained in tapping mode.
Figure 12. The average dimension of pitting in HFP and NC4 from AFM images: (a) pitting diameter and (b) pitting depth of the HFP and NC4 at various fluences. The small horizontal lines represent the value for the respective pristine samples.
the XRD pattern of NC4, revealing intercalated disordered structure. Pure PVDF shows a highly intercalated and ordered nanostructure18 of the same clay content (4 wt %), but presumably due to better interaction with the copolymer (HFP), the nanocomposite (NC4) exhibits a highly disordered structure as revealed from both the X-ray pattern and the TEM micrograph. Pure HFP does not have any nanostructure, as evident from the absence of any peak in the XRD pattern. Piezoelectric β-Phase. The XRD patterns of pristine HFP and NC’s exhibiting their crystal structures are presented in Figure 3. The structure of HFP is very similar to pure PVDF, showing the characteristic peaks of R-form at 17.6°, 18.3°, 19.9°,
and 26.4° corresponding to (020), (110), (111), and (120)/(021) crystal planes.5,6 The XRD patterns of NC’s have clearly different peaks, which are the characteristic crystal planes of β-phase at 20.1° corresponding to (200/110) planes. Hence, HFP has crystallized in piezoelectric β-form in the presence of modified nanoclay, predominantly on the surface of the nanoclays because of suitable interactions. In other words, the clay platelets are sandwiched by two crystal sheets (β -phase: all trans conformation), which in turn provides a Bragg’s reflection corresponding to the d-spacing around 1.5 nm (the thickness of nanoclayis ∼1 nm), matches with the (002) plane of the layered intercalated structure. The higher coherency of the (002) reflection, as compared to the (001) plane in Figure 1, is essentially a combined effect of higher-order reflection along with the manifestation from the sandwiched structure. The XRD pattern of low-temperature, solution-cast film also shows a similar behavior and, thereby, eliminates the possibility of collapsing of the galleries, which may occur at high-temperature, melt-processed nanocomposites. Figure 4 shows FTIR spectra for the phase structure of HFP and NC’s in the range of 1000-400 cm-1. The R peaks at 490, 615, 763, and 976 cm-1 have been clearly observed for pure HFP. A complete disappearance of those peaks has been noticed for NC4 and NC8. In contrast, strong β peaks at 510 and 840 cm-1 are evident.30,31 The absorption band at 763 cm-1 is related to in-plane bending vibration. The band at 615 cm-1 has been assigned to a mixed mode of CF2 bending and C-C-C skeletal vibration; this peak is parallel to the chain axis. The 490 cm-1 band is related to bending and wagging vibrations of the CF2 group ascribed to the R-HFP polymorph. In the case of nanocomposites, both NC4 and NC8, the band at 840 cm-1 has been assigned to a mixed mode of CH2 rocking and CF2 asymmetric stretching vibration; this mode is parallel to the chain axis. The 510 cm-1 band has been assigned to the CF2 bending mode, and this mode is perpendicular to the polar b-axis. However, FTIR studies also support the transformation of R-form to the piezoelectric β-form in the presence of nanoclay. It is worth mentioning here that the peaks from X-ray and FTIR assigned for β-phase are quite broad, indicating the presence of some mixed R-phase. The polymer chains away from the clay tactoids may crystallize in the R-phase while the polymer chains nearby clay tactoids crystallize in the β-phase. The single melting temperature of NC’s around 164 °C against the double endotherms of pure HFP (∼152 and 159 °C) reveals the transformation of the β-phase in the presence of nanoclay (Figure 5a).32 The double melting endotherms are
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Figure 13. SEM images of HFP and NC4 surfaces before and after the indicated ion irradiation. The number “0” indicates zero fluence (before irradiation).
Figure 14. The weight loss as a function of time in the swelling experiment for pure HFP and NC4 exposed to the indicated fluence. The negative values represent a loss in weight.
presumably due to the melt-recrystallization and imperfect crystallites present in pure HFP, whereas meshlike crystallites exhibit a single endotherm in NC’s. The heat of fusion gradually varies from 33 J g-1 of pure HFP against 25 J g-1 for NC8. The higher crystallization temperature (138.4 °C) for NC8, as compared to pristine HFP (123.3 °C), suggests that nanoclay acts as a nucleating agent (Figure 5b). The crystallization temperature increases with increasing nanoclay content in the NC’s. Further, the absence of any birefringence in nanocomposites against the usual spherulitic patterns in pure HFP under polarizied light (Figure 6) has also implied the phase transformation of HFP from R to β in the presence of nanoclay.33 The lack of birefringence in NC’s is due to the random orientation
of the smaller meshlike β crystallites16 against the larger, highly ordered spherulitic structure observed in pure copolymer. In NC’s, the R spherulites have not been observed, even for hightemperature (140 °C) crystallized specimens. Instead, welldispersed agglomerates of the clay particles are noticed in the HFP matrix. Thus, a distinct transition of the crystalline phase (R to β) has been observed in the presence of a few weight percentage of nanoclay through XRD, FTIR, DSC, and POM. Dynamic Mechanical Behavior. Figure 7 shows the dynamic mechanical behavior of HFP and its nanocomposites in the temperature range of -80 to 150 °C measured at 1 Hz. The nanocomposites exhibit a higher storage modulus as well as a loss modulus, as compared to pure HFP. The increment in the storage modulus varies on clay loading and the temperature zone and is presented in Table 1. The storage modulus increased by 13 and 35% for NC4 and NC8, respectively, at room temperature. However, at high temperature (∼140 °C), well above the glass transition temperature, an increment up to 75% is noted, presumably due to favorable interaction and intercalation of polymer chains inside the organically modified silicate layers. It is also observed that loss moduli of NC’s are higher than that of pure HFP, except for a small temperature window, suggesting more ductile behavior of the nanocomposites as compared to the pristine polymer. In addition, three transitions, which are more prominent in tan δ curves in Figure 7c, are evident in the temperatures ranges studied. The lower temperature transition around -40 °C occurs due to glass transition (Tg) in the amorphous phase34 and arises from the cooperative segmental motions within the main chain.35
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a The red spheres are the free radicals generated, and the green circles indicate the region of recombination of free radical leading to crosslinking.
Figure 15. (a) Gel fraction and (b) solubility parameter of pure HFP and its nanocomposite (NC4) in the swelling experiment as a function of fluence.
Figure 16. GPC traces of pure HFP and NC4 at different indicated fluences as a function of elution time. The y-axis indicates the intensities.
Further, the Tg gradually shifts (peak of tan δ curve) toward a higher temperature from -42 to -36 °C for pure HFP and NC8, respectively, due to restriction in the segmental motion from the intercalation of polymer inside galleries. The second transition, corresponding to the relaxation of crystal-amorphous and nanoclay-amorphous interphase, is observed for both pure HFP and its nanocomposites at 10 °C. This transition is evident in R-HFP34 due to the greater size of the interface in the presence of spherulites present in the R-phase, but its absence in unfilled β-PVDF36 might be due to the minimum interphase area existing between the meshlike β-crystallites and the amorphous chains. In contrast, the highly intense transition at 10 °C in nanocomposites as compared to
pure HFP is presumably due to larger interfaces between nanoclay and polymer chains. The third peak is noticed at 75 °C and can be attributed to the molecular motions in the interior of the crystalline regions, coupled with the motion of the lamellar surfaces (premelting phenomena).37 Two models have been proposed by McBrierty et al.38 to interpret (a) a rotation of crystalline chains in the vicinity of defects, and, (b) rotational oscillations of restricted amplitude of all the crystalline chains about the main axis. The third transition peaks for nanocomposites have drastically and gradually been reduced as compared to pure HFP, mainly due to the ease of rotation of smaller meshlike crystalline chains, as opposed to the bigger-size spherulites in pure HFP. Consequently, the third transition in NC’s takes place at a few degrees lower temperature (low thermal energy) vis-a`-vis pure polymer. All three transitions had also been shown in unfilled β-PVDF through dielectric relaxation and NMR techniques.35,38 The Effect of Swift Heavy Ions on Structure. The effect of SHI irradiation on HFP and its nanocomposites has been studied at a wide range of fluence from 1 × 1010 to 5 × 1012 ions/cm2. XRD patterns of HFP and NC4 before and after SHI irradiation are shown in Figure 8. The R and β crystalline structures of both pure HFP and NC4, respectively, remain unaffected after irradiation. The peak intensities gradually decrease for pristine HFP while the peak intensities have not been reduced that much for NC4. The comparative and quantitative reduction in intensities of same thickness samples for both cases is presented in the Supporting Information, Figure S1. A similar tendency of the melting behavior, before and after irradiation, of the individual crystalline phases of pure HFP and NC4 is presented in Figure 9. The heat of fusion (∆Hm) of pure HFP decreases considerably with increasing fluence, indicating less crystallinity or degradation (or both) in pure HFP, whereas a minimum decrease in the ∆Hm has been noticed for NC4 (Figure 10). The melting temperatures were almost the same for HFP (159 °C) and NC4 (163 °C) both before and after SHI
HFP Layered Silicate Nanocomposites irradiation, suggesting no change of crystalline structure. XRD and DSC measurements exhibit significantly less crystallinity or degradation in pure HFP, whereas the ceramic nanoclay and associated β-phase restrict the degradation of the matrix polymer in NC’s upon high-energy SHI irradiation. However, HFP is affected marginally in the presence of a few weight percentages of organically modified nanoclay upon SHI irradiation. A similar nature of retention of the structure but reduction in the intensity has also been found in FTIR analyses of irradiated HFP and NC’s (Supporting Information, Figure S2). There is no significant change in crystallization temperature (Tc) in HFP and NC’s after irradiation (Supporting Information, Figure S3), but nanoclay exhibits nucleating behavior (higher Tc) in NC’s. The enthalpy of crystallization gradually decreases with increasing fluence, which complementarily proves that degradation also occurs along with increasing amorphous content upon SHI bombardment. On the other hand, degradation is somehow restricted in NC4 with respect to pure HFP. The heat of fusion from the second heating indicates that ∆Hm gradually decreases with increasing fluence, which supports the degradation of HFP after SHI irradiation, although the relative degradation is quite low in the case of NC’s in comparison to pure HFP (Supporting Information, Figure S4). The Effect of SHI on the Surface Morphology. The surface morphology of HFP and NC’s has been explored by using atomic force and scanning electron microscopy. Figure 11 shows a comparison of 3-D AFM images of a 5 × 5 µm2 area of pure HFP and NC4 exposed with various fluences. Before irradiation, the surface of NC’s is smoother than that of pure HFP. The nanoclay helps to smooth the surface of the HFP; however, SHI irradiation causes huge surface roughening in pristine HFP, and the damage is enhanced with increasing fluence while the surface roughness does not change much after irradiation in NC4, indicating radiation-resistant behavior of the nanocomposites. The dimensions of the pittings on the surfaces caused by irradiation have been shown quantitatively in Figure 12. Both the diameter and depth of pittings have increased significantly with fluence in HFP, which clearly suggests huge damage on the surface with a percolation threshold of ∼5 × 1011 fluence. In contrast, the pitting dimensions for NC4 change a little (almost parallel to x-axis) with increasing fluence. NC4 shows a 320 nm average pitting diameter and 24 nm pitting depth at 5 × 1012 fluence, in contrast to a 1500 nm pitting diameter and 200 nm depths for pure HFP in the same fluence. In fact, pure HFP film turns slightly black and brittle after being exposed to 1 × 1012 fluence, but the NC’s are brown and ductile even at higher fluence (up to 5 × 1012). The surface morphology through SEM also indicates that severe damage occurs in pure HFP after SHI irradiation, whereas a smooth surface prevails in NC4 after the same dose of irradiation (Figure 13). The ceramic nanoclay with its ultralarge surface area uniformly distributed in the matrix along with tiny, compact β-crystallites protect the nanocomposites surface from damage by high-energy, swift, heavy ions. The Effect of SHI on Gelation. High-energy, swift, heavy ions are known to induce chain session, resulting in free radicals and evolved gas. The free radicals may recombine, leading to the formation of cross-linking between polymer chains. The cross-linked part is usually insoluble in common solvents of the polymer. Both pristine HFP and NC’s are readily soluble in hot DMF, but after irradiation at high fluence, discrete films have been observed in hot solvent, even after 48 h of extraction. A part of HFP has been solubilized in swelling studies, and weight loss is shown in Figure 14. It clearly shows that irradiated
J. Phys. Chem. B, Vol. 113, No. 34, 2009 11639 nanocomposites lose minimum weight, as compared to pristine HFP under the same conditions. To understand the extent of the cross-linking or gel fraction, a solubility test has been conducted, and the gel fraction has been plotted as a function of fluence in Figure 15a. The gel fraction increases with fluence, but the relative rate is considerably higher for NC4 as compared to pure HFP, especially beyond 1011 fluence. This indicates that HFP is cross-linked at a greater extent in the presence of nanoclay in NC’s. For comparing the chain session and cross-linking, the classical Charlesby-Pinner equation39,40 has been employed:
S + √S )
Gs A + 2Gc D
(2)
where S ) 1 - F is the soluble fraction, Gs and Gc are the radiation-chemical yields of chain session and cross-linking, respectively, D is the fluence, and A is a constant. The evolution of S + Sj vs 1/D is plotted in Figure 15b. A higher solubility fraction is evident for pure HFP, as compared to NC4. The ratios of Gs/Gc have been calculated from the extrapolation of the linear part of the curves at 1/D ) 0 (D f∝) as 3.5 and 3.0 for pure HFP and NC4, respectively. The higher value of Gs/Gc for pure HFP suggests greater chain session and less crosslinking. On the other hand, cross-linking is significant in NC4. If the cross-linking occurs between polymer chains (interchain cross-linking), the molecular weight should increase. The effect of the molecular weight on the SHI irradiation is compared in Figure 16 for both pure HFP and NC4. The molecular weight distribution is bimodal before irradiation. The intensity of the main peak at longer retention time gradually decreases with increasing fluence for both pure HFP and NC4, but certainly has a greater impact on NC4. In addition, an extra peak at ∼29.5 min retention time has appeared for SHI-irradiated samples, and its intensity gradually increases with increasing fluence. The molecular weight, corresponding to the peak position (29.5 min retention time), as calculated from the extended calibration curve shown in Supporting Information Figure S5 is 44 00 000 against the average molecular weight j n) of pristine HFP of 2 00 000. Further, the peak intensity (M becomes more prominent, even at low fluence, for NC4 than for that of pure HFP, suggesting a greater extent of cross-linking phenomena in the presence of nanoclay in the NC’s. It is worth mentioning here that molecular weight measurement was performed using 48 h extracted samples. A highly cross-linked sample was intact as a film in the solution. However, combining the molecular weight measurement and solubility experiment, we conclude that chain session is significant in pure HFP, but the extent of cross-linking is greater in nanocomposites, even though both the processes occurred simultaneously in pure HFP as well as in nanocomposites. It is clear from the above discussion that the extent of crosslinking is greater in nanocomposites, as compared to pure HFP; that is, nanoclay helps in cross-linking or recombining free radicals. On the other hand, degradation/chain session occurs in pure HFP upon SHI irradiation. If a high-energy ion bombards a polymer matrix, it loses its energy either by interacting with target nuclei, called nuclear stopping, or by interacting with a target electron, called electronic stopping. Electronic stopping is largely responsible for cross-linking, and nuclear stopping is for chain session, although both processes may cause crosslinking as well as chain session.4 A number of chemical species, which includes exciton phonon, ion pair, and radical pair, usually
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generate upon SHI bombardment. The formation of a chemical bond by connecting two freely dangling ions or radical pairs on neighboring molecules causes cross-linking. It has already been shown that the HFP molecules crystallized in all transconfiguration β-phase on the surface of the layered silicates where parallel molecular orientation prevails. Once free radicals formed in the HFP molecule by SHI bombardment, due to its confinement with layered silicate and close proximity of another free radical in the neighboring straight chain of β-phase, the probability of recombination of the free radical gets facilitated in nanocomposites. In contrast, the probability of recombination of free radicals in R-phase (zig-zag) conformation in pure HFP is quite less due to a very low time width (10-15 s) over which an absorption of energy can take place. The time was calculated from the uncertainty principle, ∆t∆E ≈ p for a relativistic particle (∼100 MeV/nucleon, similar to our measurement).4 It is worth mentioning here that molecular motion and diffusion processes start at ∼10-13 and 10-12 s, respectively.41 Therefore, the chances of recombination of free radicals are much less in pure HFP, whereas the parallel and restricted molecular structure in the presence of layered silicates favors the probability of recombination in β-HFP in nanocomposites. The formation of free radicals in polymer chains with and without the presence of layered silicate is presented in Scheme I. The probability of the formation of a cross-link is greater because the free radicals reside in the parallel straight chains of the β-phase in nanocomposites (shown by the green circles). The above experiments demonstrate that a few weight percentage of layered silicate dispersed in HFP makes the matrix more flexible and can be used as radiation-resistant materials. The same idea can also be extended to other thermoplastics or engineering plastics. Other potential uses of irradiated HFP nanocomposites can be for fuel cell membranes in which both the strength and right kind of porosities are required. It is also possible to generate ion tracks of different dimensions by irradiating SHI into a polymeric substrate, followed by suitable etching. We are continuing work in that direction, as well. Conclusions The nanocomposites of HFP copolymer with organically layered silicate have been prepared through a melt extrusion technique. The intercalated and partially exfoliated nanostructures have been revealed depending on the amount of nanoclay present in the matrix. The metastable piezoelectric β-phase got stabilized in the presence of nanoclay platelets, which act as nucleating agents, as well. Both the storage and loss modulus of HFP increase with increasing nanoclay content, suggesting the excellent reinforcing nature of nanoclay, and at the same time, the flexibility of HFP is also enhanced in nanocomposites. The different transitions in tan δ curves have been interpreted on the basis of the structural aspects of the HFP with and without the presence of disk-like nanoclay. The piezoelectric β-phase is retained in nanocomposites after SHI irradiation, indicating that the nanocomposites can be used as radiation-resistant materials at high temperature. Amorphous content is enhanced, and degradation takes place in pure HFP after SHI bombardment, whereas a slight change is observed in nanocomposites in the same condition. SHI induces pittings in pure HFP, and its dimensions and degradation are enhanced significantly beyond 5 × 1011 ions/cm2 fluence, which limit its use for any ion irradiation applications. In contrast, the nanocomposites exhibit almost no change in the whole range of fluence studied. The swelling experiments and gel fraction studies of irradiated samples indicate that chain session is a
Tiwari et al. major occurrence in pure HFP, whereas cross-linking is a core phenomenon in nanocomposites, presumably due to the parallel chain configuration of HFP molecules on the surface of nanoclay discs. The molecular weight measurement and residual weight also support the enhanced cross-linking trend with increasing fluence. Acknowledgment. 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. Asish Lele, Dr. S. Sinha Ray, Dr. B. Ray, and Prof. A. S. K. Sinha for extrusion work, TEM, GPC, and FTIR studies, respectively. The kind supply of HFP samples by Ausimont, Italy, is highly acknowledged. Supporting Information Available: Additional information as noted in text. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Lyons, B. J. Radiat. Phys. Chem. 1995, 45 (2), 159. (2) Biswas, A.; Gupta, R.; Kumar, N.; Avasthi, D. K.; Singh, J. P.; Lotha, S.; Fink, D.; Paul, S. N.; Bose, S. K. Appl. Phys. Lett. 2001, 78 (26), 4136. (3) Percolla, R.; Calcagno, L.; Foti, G.; Ciavola, G. Appl. Phys. Lett. 1994, 65 (23), 2966. (4) Lee, E. H. Nucl. Instrum. Methods Phys. Res., Sect. B 1999, 151, 29. (5) Huan, Y.; Liu, Y.; Yang, Y. Polym. Eng. Sci. 2007, 47 (10), 1630. (6) He, X.; Yao, K.; Gan, B. K. J. Appl. Phys. 2005, 97, 084101. (7) Moggi, G.; Bonardelli, P.; Bart, C. J. Polym. Bull. 1982, 7, 115. (8) Adem, E.; Rickards, J.; Burillo, G.; Avalos-Borja, M. Radiat. Phys. Chem. 1999, 54, 637. (9) Tashiro, K.; Tadokoro, H.; Kobayashi, M. Ferroelectrics 1981, 32, 167. (10) McGrath, J. C.; Ward, I. M. Polymer 1980, 21, 855. Tamura, M.; Ogasawara, K.; Ono, N.; Hagiwara, S. J. Appl. Phys. 1974, 45, 3768. (11) Scheinbeim, J.; Nakafuku, C.; Newman, B. A.; Pae, K. D. J. Appl. Phys. 1979, 50, 4399. Newman, B. A.; Yoon, C. H.; Pae, K. d.; Scheinbeim, J. J. Appl. Phys. 1978, 49, 4601. (12) Wang, J.; Li, H.; Liu, J.; Duan, Y.; Jiang, S.; Yan, S. J. Am. Chem. Soc. 2003, 125, 1496. (13) Lovinger, A. J. Polymer 1981, 22, 412. (14) Priya, L.; Jog, J. P. J. Polym. Sci., Part B: Polym. Phys. 2002, 40, 1682. (15) Priya, L.; Jog, J. P. J. Polym. Sci., Part B: Polym. Phys. 2003, 41, 31. (16) Shah, D.; Maiti, P.; Gunn, E.; Schmidt, D. F.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. AdV. Mater. 2004, 16 (14), 1173. (17) Shah, D.; Maiti, P.; Jiang, D. D.; Batt, C. A.; Giannelis, E. P. AdV. Mater. 2005, 17 (5), 525. (18) Tiwari, V. K.; Kulriya, P. K.; Avasthi, D. K.; Maiti, P. ACS Appl. Mater. Interfaces 2009, 1 (2), 311. (19) Zhudi, Z.; Wenxue, Y.; Xinfang, C. Radiat. Phys. Chem. 2002, 65, 173. (20) Pae, K. D.; Bhateja; Gilbert, J. R. J. Polym. Sci., Part B: Polym. Phys. 1987, 25, 717. (21) Chailley, V.; Balanzat, E.; Dooryhee, E. Nucl. Instrum. Methods Phys. Res., Sect. B B 1995, 105, 110. (22) Torrisi, L.; Percolla, R. Nucl. Instrum. Methods Phys. Res., Sect. B 1996, 117, 387. (23) Szenes, G. Nucl. Instrum. Methods Phys. Res., Sect. B 1999, 155, 301. (24) Mazzei, R. O.; Bermudez, G. G.; Tadey, D.; Rocco, C. Nucl. Instrum. Methods Phys. Res., Sect. B 2004, 218, 313. (25) Soresi, B.; Quartarone, E.; Mustarelli, P.; Magistris, A.; Chiodelli, G. Solid State Ionics 2004, 166, 383. (26) Kulshrestha, V.; Awasthi, K.; Vijay, Y. K. Int. J. Hydrogen Energy 2007, 32, 3105. (27) Laghari, J. R.; Hammoud, A. N. IEEE Trans. Nucl. Sci. 1990, 37 (2), 1076. (28) Bersack, J. P.; Haggmark, L. G. Nucl. Instrum. Methods 1980, 174, 257. (29) Maiti, P.; Nam, P. H.; Okamoto, M.; Hasegawa, N.; Usuki, A. Macromolecules 2002, 35, 2042. Maiti, P.; Yamada, K.; Okamoto, M.; Ueda, K.; Okamoto, K. Chem. Mater. 2002, 14, 4654.
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