Article pubs.acs.org/JPCB
Free Volume Study on the Origin of Dielectric Constant in a FluorineContaining Polyimide Blend: Poly(vinylidene fluoride-co-hexafluoro propylene)/Poly(ether imide) R. Ramani,† V. Das,† A. Singh,† R. Ramachandran,‡ G. Amarendra,‡ and S. Alam*,† †
Polymer Science Division, Defence Materials and Stores Research and Development Establishment, G.T. Road, Kanpur 208 013, India ‡ Materials Physics Division, Materials Science Group, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, India S Supporting Information *
ABSTRACT: The dielectric constant of fluorinated polymides, their blends, and composites is known to decrease with the increase in free volume due to a decrease in the number of polarizable groups per unit volume. Herein, we report an interesting finding on the origin of dielectric constant in a polymer blend prepared using a fluorine-containing polymer and a polyimide probed in terms of its available free volume, which is distinct from the generally observed behavior in fluorinated polyimides. For this study, a blend of poly(vinylidene fluoride-co-hexafluoro propylene) and poly(ether imide) was chosen and the interaction between them was studied using FTIR, XRD, TGA, and SEM. The blend was investigated by positron annihilation lifetime spectroscopy (PALS), Doppler broadening (DB), and dielectric analysis (DEA). With the increase in the free volume content in the blend, surprisingly, the dielectric constant also increases and is attributed to additional space available for the polarizable groups to orient themselves to the applied electric field. The results obtained would pave the way for more effective design of polymeric electrical charge storage devices. and phenyl, since they hinder efficient chain packing.7 This high degree of free volume decreases the number of polarizable groups/unit volume and, thus, limits the dielectric constant.6,8 Apart from the arrangement of constituent atoms and free volume of the polymer, the moisture content in a polymer can also influence the dielectric constant. Addition of hydrophobic atoms like fluorine into a polymer can diminish the moisture content and decrease the dielectric constant.6,7 Such fluorine containing polymers are grouped as fluoropolymers and are attractive niche macromolecules because of their versatility and ability to display a unique combination of relevant properties.9 Among fluoropolymers, poly(vinylidene fluoride) (PVDF) and its copolymers are technologically important polymers possessing high dielectric constant and are available in various crystalline forms.9,10 Because of its high dielectric constant, PVDF has been considered as an ideal candidate for high charge storage capacitor applications.10 The incorporation of amorphous hexafluoropropylene (HFP) into the VDF main constituent significantly reduces the crystallinity of PVDF.10,11 The amorphousity in this copolymer arises mainly from the pendant trifluoromethyl group (−CF3) of HFP
1. INTRODUCTION Polymer blending offers an attractive and simpler method in the development of new polymeric materials compared to chemical synthesis, and the blends often exhibit more desirable characteristics than individual polymers.1,2 The blending allows optimization of some of the properties of the homopolymers and provides an economic way to create new materials with desired properties. Thus, blending of polymers has been an important industrial approach toward the development of novel polymeric materials.1,3 When a polymer blend is formed, free volume is one of the properties that gets affected4,5 and the free volume is closely related to dielectric constant (ε′).6,7 The relationship between these two physical properties is the nucleus of this study. The dielectric constant of a polymer is mainly influenced by two parameters. First is the molecular polarizability that is decided by the constituent atoms that have an appropriate electronegativity and bonding in them in a suitable configuration. Modifying the type and number of polarizable groups, the dielectric constant of a polymer can be changed.6 Second is the free volume associated with the polymer. Since the dielectric constant of air is close to one, introducing free volume (porosity) to a polymer reduces the dielectric constant of the polymer.6 The free volume in a polymer can be modified by incorporating space occupying groups such as ethyl, methyl, © 2014 American Chemical Society
Received: June 17, 2014 Revised: August 15, 2014 Published: September 30, 2014 12282
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Figure 1. Chemical formulas of the polymers used in this work and the possible interaction site between them in their blend.
to reduce the interchain interactions, prohibits close packing of chains, and decreases the ability of charge transfer complexes (CTCs) between dianhydride and diamine moieties and thus serves as molecular spacers leading to increase of free volume in them.13,21 However, the fluoro-polyimides have drawbacks of low mechanical strength and high thermal expansion coefficient.23 Poly(ether imide) (PEI) is a high performance aromatic amorphous thermoplastic polyimide aimed at replacing metals as well as targeting other specialty applications such as electronics, automotives, and membranes.24 PEI has useful properties such as superior thermal stability, high modulus and toughness.14,24 Incorporation of fluorine atoms into PEI backbone by the attachment of −CF3 groups was reported to produce low dielectric properties compared to PEI.13,25 To the best of our knowledge, the effect of fluorine atoms on the dielectric constant and free volume properties of a polyimide has not been studied by blending it with a fluorine-containing polymer. Previous studies have only focused on understanding the dielectric constant and free volume by altering the polyimide structure using different fluorinecontaining dianhydride and diamine moieties.6,8,16,19,20,25−27 The PVDF is compatible with polymers having a >CO group in its structure.2,9,28−30 Thus, P(VDF-co-HFP) and PEI (having four carbonyl groups in its structure; see Figure 1) are expected to be compatible. Additionally, this copolymer has a pendant group (−CF3) in its structure. Thus, these two polymers are chosen for the study and their blends were made in various proportions. The type of interaction in them, the free volume, rheological, dielectric, and thermal and viscoelastic properties of the blend were studied. Interestingly, our results show increase in the dielectric constant with increase in the free volume, a result different from that generally reported in polyimides,6,17−19 their blends,31,32 and more recently even in polyimide block copolymer aerogels.33 The reason for this positive deviation of dielectric constant with free volume has been identified. For evaluating the free volume, we have used the well established positron annihilation lifetime spectroscopy (PALS), a highly sensitive and nondestructive tool that can quantitatively measure the size of free volume in polymers and polymer blends.4,5,34−37 In the physical aspect, PALS considers the two most fundamental parameters time and dimension. This technique is based on the fact that the lifetime of positron and its bound state with an electron called positronium (Ps) are sensitive to the existence of structural inhomogeneities in polymers.34 This localized annihilation in “free volume holes” makes it the microprobe of choice since its lifetime becomes a
that does not allow efficient chain packing and thus contributes to free volume.8 The copolymer P(VDF-co-HFP) finds application as a polymer electrolyte since the amorphous HFP permits diffusion of liquid electrolytes and the crystalline regions of VDF help to maintain the mechanical integrity of the film.12 Polyimides are another set of polymers that are rigid, thermally stable, and obtained by polycondensation reactions of diamines with dianhydrides.13 Apart from thermal stability, polyimides are endowed with high glass transition temperature (Tg), good mechanical and electrical properties, low dielectric constant, high breakdown voltage, lightweight, flexibility, and inertness to solvent.13 They find extensive use in aerospace and electronic applications.14 More information on advanced polyimides and their applications can be found in a recent review.15 Various strategies have been adopted in the past to decrease the dielectric constant of polymides like incorporating suitable diamine and dianhydride moieties (that minimizes polarizability and provides high degree of free volume) as well as adding fluorine atoms into the polyimide molecular structure.6,8,16 Fluorine replacement for hydrogen is known to decrease the local electronic polarization since the electronic polarizability of the C−F bond is less than that of the C−H bond.6,7 An important contribution in understanding the molecular process that lead to dielectric properties in fluorine-containing polyimides was made at NASA-Langley.17,18 The authors have studied the effect of change in free volume on the dielectric constant in them and found a correlation of high free volume content and low dielectric constant.17,18 Later, other groups have also found that the ε′ of polyimide decreases with increasing wt % of fluorine in different polyimides.6,8,16,19,20 This is because, fluorine substitution to polyimides increases the free volume due to greater stearic volume of fluorine relative to hydrogen which inhibits efficient chain packing and thus decreases the dielectric constant.6,8 As explained above, since the introduction of fluorine into polyimides brings about dramatic improvements in many properties of polyimides, it has become a subject of intense research and review.13 Such fluorine-containing polyimides form an interesting group since, apart from their low dielectric constant, they also exhibit low water uptake and possess relatively high permeability as compared to their fluorine-free counterparts.6,16,21,22 Also, the introduction of −CF3 groups into polyimide has been one of the most widely used strategies for structural modification leading to substantial property enhancement.13,21 This is because, in comparison to fluorinefree polyimides, the −CF3 group in the polyimide chain tends 12283
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ments were carried out in dynamic conditions at a heating rate of 20 °C/min both in nitrogen and in dry air atmosphere (flow rate = 50 mL/min). 2.2.4. Scanning Electron Microscopy (SEM). The surface morphology of the blends was examined using a Carl Zeiss EVO 50 SEM, with 10 kV operating voltage. The fracture surfaces of the samples were obtained by quenching the samples in liquid nitrogen and breaking them. The blends were then treated with acetone at room temperature for 24 h, in order to remove PVH phase from them. The acetone etched blends were thoroughly dried and were sputter coated with gold on the viewing surface to enhance its conductivity. 2.2.5. Positron Annihilation Lifetime Spectroscopy (PALS) and Doppler Broadening (DB) Measurements. In PALS method, positrons are injected into polymers by a radioactive source. When positrons enter into polymers, a fraction of them form positronium (Ps), which is a bound state of electron and positron. Depending on the spin of electron and positron, both para Ps (p-Ps) and ortho Ps (o-Ps) are formed. In p-Ps, the spin of electron and positron are antiparallel and has an intrinsic lifetime of 0.125 ns, whereas in o-Ps the spins of both the particles are parallel and has an intrinsic lifetime of 140 ns. In polymers, the o-Ps is formed in free volume sites, where the positron of o-Ps picks up an electron of opposite spin at the free volume cavity wall for annihilation. This process is called pickoff annihilation by which the o-Ps lifetime gets reduced to 1−5 ns. The pick-off annihilation lifetime is inversely proportional to the overlap of the positron and electron wave function and thus can be related to the size of the “free volume hole”, where the annihilation takes place.34 Both o-Ps pick-off and p-Ps intrinsic annihilation results to two γ photon emission with average energy centered around 511 keV. In the Doppler broadening (DB) method, of the two γ-photons produced upon annihilation, the energy of one is upshifted while the other is downshifted by an amount given by ΔE = PLc/2, due to Doppler effect. Here, PL is the longitudinal component of momentum of the pair along the direction of γray emission and c is the velocity of light. Since the positron is thermalized in the sample, the momentum comes from the energetic electrons.38,39 The DB spectrum can be analyzed using S and W parameters using which the information regarding the defect can be deduced. The S-parameter is defined as the ratio of area under central part of the spectrum to the total area, while the W-parameter is the ratio of sum of the area under the wings in a fixed window over the total area of the spectrum.38,41,42 High S value for a given material suggests more annihilation with low-momentum electrons, while high W value suggests more annihilation with highmomentum electrons. More details on PALS and DB methods can be found from the literature.34,41,42 The positron annihilation lifetime measurements were performed at room temperature using a conventional fast− fast coincidence system having a time resolution of about 200 ps. A positron source (22Na) of 10 μCi was sandwiched between two pieces of identical samples. The lifetime spectra were resolved into three lifetime components using PATFIT program43 after subtracting the source and background corrections. The three component analysis gave better convergence than two and four component fits. The positron lifetime spectra were also analyzed using the MELT program44 to obtain a continuous distribution of the annihilation lifetime. The DB measurement system consists of ORTEC highpurity germanium (HpGe) detector having 1.3 keV energy
measure of the electron density and free volume size, both of which helps to understand the polymer microstructure.34 The change in free volume size together with its distribution pattern plays an imperative role in understanding the miscibility aspect of polymer blends.4,5 In addition to the average free volume size and its distribution, knowledge on the chemical environment inside the free volume holes can be obtained from the Doppler broadening (DB) of the annihilating radiation.35−39 In the DB method, the average energy shift from the 511 keV of the two γ-photons that are produced by the annihilation of electronpositron pair is measured.34,38 The interesting part of the positron results is that the ortho-positronium (o-Ps) intensity (I3) increases with increase in fluorine content in the blend; while o-Ps inhibition process leading to decrease in I3 is commonly reported in halogenated polymers and their blends.35,38,40 In the present study, we have used PALS together with DB to study the free volume in the blend and the dielectric constant was measured using a dielectric analyzer. Additionally, the type of interaction between P(VDF-co-HFP) and PEI in their blend and their rheological, thermal, and viscoelastic behaviors are discussed.
2. EXPERIMENTAL SECTION 2.1. Sample Preparation. The copolymer of vinylidene fluoride and hexa fluoropropylene P(VDF-co-HFP) having Mw = 400 000 g/mol was obtained from Sigma-Aldrich in pellet form. Henceforth, this copolymer will be termed as PVH. The PEI of grade Ultem-1000 with Mw = 55 000 g/mol was obtained in powder form from SABIC innovative plastics, Europe. The PVH was dried in vacuum at 50 °C for 24 h, while PEI was vacuum-dried for 48 h at 120 °C before use. Blends with various weight percentage of PVH/PEI were prepared and designated as PHPI00, PHPI20, PHPI40, PHPI50, PHPI60, PHPI80, PHPI85, PHPI90, PHPI95 and PHPI100, where the subscripts denote the wt % of PEI in the blend. The blending was performed using a twin screw extruder (make Thermo Scientific, Germany; model Haake Minilab II) at 350 °C. The strand leaving the extruder was quenched, air-dried, and chopped into granulates. These were subsequently injection molded at 370 °C with an injection molding machine (make Thermo Scientific, Germany; model Haake Minijet II) into rectangular bars of size 50 × 10 × 3 mm3. The chemical structures of PVH and PEI and the possible interaction site between them in their blend are shown in Figure 1. 2.2. Characterization. 2.2.1. Fourier Transform Infrared Analysis. The FTIR spectrophotometer of PerkinElmer (model Spectrum 100) was used to identify the possible interacting groups between the polymers. The samples were ground with KBr powder using a pestle and mortar and pressed into pellets. The samples were then dried in the vacuum oven at least for 12 h before the infrared spectra were recorded. 2.2.2. X-ray Diffraction (XRD). The XRD measurements were performed with PAN Analytical X-ray systems (model X’pert PRO) (Cu Kα radiation, λ = 1.54 Å) in the 2θ range of 5° to 80°. A scanning rate of 0.05° s−1 was used to record the XRD pattern. During the X-ray scattering experiments, the samples were sealed between two thin Kapton foils and the raw intensity data were corrected for background scattering. 2.2.3. Thermogravimetric Analysis (TGA). Thermo gravimetric analysis was performed using a TGA instrument (Mettler Toledo TGA/SDTA 851e) in the temperature range 30 to 1000 °C. Samples having mass ca. 5 mg were placed in 100 μL aluminum crucibles. Thermal decomposition experi12284
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Figure 2. FTIR spectra for the PVH/PEI blend in the (A) 800−940 cm−1 region, (B) 1200−1450 cm−1 region, and (C) 1650−1850 cm−1 region. The spectra are up shifted for clarity.
resolution at 662 keV gamma line of 137Cs source. The sourcesample sandwich was placed at a distance of 10 cm from the detector and the data was collected to obtain 1 million counts with an average time for each run to be ca. 1.5 h. 2.2.6. Rheological Measurements. The rheological measurements were performed using a Rosand capillary rheometer of Malvern make (model RH7) at 380 °C in air medium. A capillary die with an L/D ratio of 16 and a diameter of 1 mm was used and the shear viscosity was measured in the shear rate range 100−1000 s−1. The shear viscosity (η) and shear stress (τs) were measured in melt state as a function of shear rate (dγ/ dt). 2.2.7. Dielectric Measurements. The dielectric measurements were conducted in the temperature range of −100 to 250 °C (up to 120 °C only when the PEI content is CH2)9,10 or can undergo dipolar interaction through the >CF2 group.47 Thus, PVH is also expected to show interaction through these two modes. Figure 2 (A-C) reveals the FTIR spectra for PVH, PEI and their blends in the 800−940, 1200− 1450, and 1650−1850 cm−1 regions, respectively. 12285
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reported, which also infers that >CF2 is an active site for dipolar type interaction52 and thus supports our results. In summary, FTIR results reveal that there is no hydrogen bonding interaction between >CO groups of PEI and >CH2 groups of PVH. Instead, a dipolar type interaction exists between the electron deficient imide group of PEI and electronegative >CF2 groups of PVH (see Figure 1). This result is really surprising in comparison with the ability of PVDF to form hydrogen bonds with polyimides.21,48 This could be speculated as being due to the presence of a −CF3 group in the HFP part of the copolymer. This speculation is based on the fact that the −CF3 group is responsible for change in polarizability of this blend8 and this would be cleared from the discussions in the subsequent sections. A careful observation of the FTIR data also reveals that the spectral features show significant changes when the PVH content is ≤20% in the blend (see Figure 2A,B). 3.2. XRD Results. The XRD profiles of the blend along with their pure polymers are presented in the 2θ range 5° to 80° (see Figure 3). The main characteristic scattering patterns
The PVDF was reported to be compatible with acrylic resins, poly(methyl methacrylate) (PMMA) and poly(ethyl methacrylate) because of hydrogen bonding between −CH2 groups of PVDF and >CO groups in these polymers.9,28,48 The recent work on a PVH/PMMA blend suggests that it contains a more miscible amorphous phase than the PVDF/PMMA blend revealing higher compatibility of PVH and PMMA.2 In our work on pyrene functionalized polyaniline (pf-PANI)/PVH blend, we have also identified hydrogen bonding interaction between the >CO groups of pf-PANI and the CH2 group of PVH.30 Furthermore, polyimides are known to form compatible blend with PVDF because of hydrogen bonding interaction with imide >CO groups.48 Hence, it is anticipated that PVH will also interact through hydrogen bonding with PEI. If there is a hydrogen bonding between >CO groups of PEI and the CH2 group of PVH, a downshift in the CO stretch of PEI is expected.30,46,49 A careful observation of the CO stretch of the imide ring near 1780 cm−1 (asym CO stretch) and 1725 cm−1 (sym CO stretch)14,46 reveals that the peak position remains stationary regardless of the blend composition (Figure 2C) indicating a lack of significant interaction at this site. Next, we analyze the 1200−1450 cm−1 region (Figure 2B) for the possible interaction at the imide (−N < ) ring of PEI. The imide ring absorbance peaks are seen at 1351 and 1393 cm−1, and the values are close to that reported.45 The peak at ca. 1393 cm−1 is due to C−N stretch of the imide ring.14,24 With the increase in PVH content, a downshift in the 1393 cm−1 peak is seen while the spectral feature at 1351 cm−1 becomes broad and shifts to higher side. These spectral changes indicate that the imide rings of PEI take part in the interaction with PVH. Based on the possible interactions mentioned earlier, the present changes could be considered as due to dipolar type interaction between the imide group of PEI and the >CF2 group of PVDF. Now, we will verify this dipolar interaction from the perspective of >CF2 groups in PVH. The asymmetric and symmetric stretching vibrations of >CF2 are reported at 1085 and 1190 cm−1, respectively.50 The overlapping spectral features of PEI pose difficulty for analysis at this region. Instead, the skeletal band of PVH at 874 cm−1 and the CH2 rocking band at 839 cm−1 are considered.47,51 More precisely, the band at 839 cm−1 is assigned to the mixed mode of >CH2 rocking and >CF2 asymmetric stretching, while the band at 874 cm−1 is due to combined >CF2 and C−C symmetric stretching vibrations.51 Please note that the peak close to 840 cm−1 and the one near 1271 cm−1 (see Figure 2B) are reported to represent the β-form of PVDF which we will discuss in the next section.10,51 The overlapping PEI peak at this region again makes it difficult to identify the spectral changes due to interaction. However, even with little addition of PVH to PEI (as low as 5%), the 874 cm−1 band slightly upshifts and becomes broad and its intensity decreases. With the increase in PVH content, this peak shifts a little more to the right and gains intensity. These spectral changes confirm that a dipolar type interaction exists between PEI and PVH. In an earlier work on PVDF/Nylon-11 blend, intermolecular interactions were observed between CONH of Nylon-11 and >CF2 group of PVDF, due to which the peak close to 880 cm−1 showed upshift and became broad in agreement with our results shown by 874 cm−1 peak.47 In another work, interaction of >CO groups in dialkyl phthalates with >CF2 group of poly(vinylidene fluorideco-chloro trifluoro ethylene) (PVDF-co-CTFE) has been
Figure 3. XRD profiles of PVH/PEI blend with various contents of PEI. The profiles are vertically shifted for clarity.
appear at 2θ < 60° and clearly reveals the change in crystalline content of PVH upon blending with PEI. The scattering profile of PVH reveals two sharp close diffraction peaks at 2θ = 18.2° (peak I) and 20.7° (peak II), a well separated peak at 26.6° (peak III), and a broad peak close to 40° (peak IV). The peaks at 18.2° and 26.6° are due to α-crystalline phase of PVDF.10,11,29 The peak at 20.7° is attributed to the β-phase of PVDF.11,30 The broad diffraction pattern seen close to 40° is also a characteristic peak of crystalline PVDF.53 The XRD profile of PEI reveals a broad amorphous halo at 2θ close to 20° and another broad hump with less intensity at ca. 43°.14,25 It is interesting to note that with the addition of PEI, even though the peak representing α-phase (2θ = 18.2°) become indistinguishable when the PEI content is ≥80 (PVH content is ≤20%), the β-peak (2θ = 20.7°) is distinctly seen. The β-form of PVDF is the thermodynamically stable form and is more polar compared to the α-form.7 The β-form has an all12286
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Figure 4. (A and B) TGA thermograms of PVH/PEI blend with varied PEI content measured in argon medium and (C and D) their respective derivative curves.
polymers, which is an indication of noncompatibility between the two polymers. Interestingly, when the PEI content reaches 80%, the first decomposition temperature (related to PVH) in the blend shows an upshift and occurs at ca. 464 °C with a decomposition maximum at ca. 504 °C (Figure 4A,C). Further increase in PEI content mildly improves the decomposition temperature of PVH, and when the PEI content is 95%, the decomposition temperature of PVH occurs at 468 °C with the decomposition maximum at ca. 508 °C. Please note that the decomposition temperature of pure PVH is 440 °C and blending with PEI improved its thermal stability by ca. 30 °C. The significant interaction between PVH and PEI is evident when the PVH content is ≤20%, which is in agreement with FTIR and XRD results. In order to further confirm the significant interaction between PVH and PEI in the PEI rich phase (when PVH ≤ 20%), the TGA measurements were also conducted in air medium (see Figure S1 (A-D) in the Supporting Information). The introduction of fluorine is known to improve the thermooxidative stability of polyimides.25 The PEI exhibits two step decomposition when pyrolyzed in air medium.54 The first one occurs at ca. 530 °C which is due to chain-scission of the backbone. The second one that has a decomposition maximum at ca. 650 °C is due to thermo-oxidation.54 Even with 5% addition of PVH (sample PHPI95), the thermo-oxidative decomposition maximum of PEI shows an upshift from 650 to 698 °C (see Figure S1(C)) and reaches to ca. 703 °C when the PVH content is ≤20% (see Figure S1(D)). The improvement in thermo-oxidation of PEI is another clear
trans conformation of the PVDF chains and is obtained when it is strained, stretched, or quenched.9 Since the samples are prepared by extrusion, the presence of the β-form is not surprising. The sharp crystalline peaks of PVH seen at 2θ = 26.6° and the other broad peak seen at higher 2θ range close to 40° become less sharp and lose their identity when the PEI content in the blend is ≥80%. This reveals that there is some level of intermolecular interactions between PVH and PEI leading to amorphous miscibility. Thus, PVH becomes disordered when the PEI content in the blend is high. Thus, XRD results clarify that PVH exists mainly in the βphase and interacts significantly with PEI when the PEI content is high (≥80%). This is in agreement with FTIR results. 3.3. TGA Results. Additional support for the interaction between PVH and PEI and the composition dependence on the extent of interaction can be obtained from TGA measurements. The TGA results for the blend samples carried out in argon medium in the temperature range 30−1000 °C are presented in Figure 4A,B, and their corresponding derivative curves are shown in Figure 4C,D. The pure PVH shows a single decomposition at a temperature of ca. 440 °C with a decomposition maximum at ca. 492 °C and a residue of ca. 4.5%.30 The high thermal stability of this polymer is due to the presence of strong C−F bonds.13 The PEI showed a decomposition pattern with no significant weight loss below 530 °C but with residue of ca. 52% and a decomposition maximum at 560 °C.14,54 With the addition of PEI to PVH, two decomposition patterns can be seen in the blend at the respective decomposition temperatures of their constituent 12287
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Figure 5. Scanning electron micrographs of the samples (A) PHPI95, (B) PHPI90, (C) PHPI80, (D PHPI60, (E) PHPI40, and (F) PHPI20.
Figure 6. Variation of (A) o-Ps lifetime (τ3) and free volume hole size (Vf3) and (B) o-Ps intensity (I3) and relative fractional free volume (Fvr) as a function of PEI content. The lines drawn are guides to eyes.
evidence that PVH interacts through fluorine group with PEI significantly when the PVH content is ≤20%. From the combined FTIR, XRD, and TGA measurements, it is clear that the significant intermolecular interaction window between PVH and PEI in their blend for the composition range studied is 5% ≤ PVH ≤ 20% or 95% ≥ PEI ≥ 80%. 3.4. SEM Results. Figure 5A−F shows the SEM photomicrographs of the PVH/PEI blend with compositions 5/95, 10/90, 20/80, 40/60, 60/40, and 80/20. These micrographs reveal the surface morphology of the blends in which PVH has been etched out by treating them with acetone. Generally, the low viscous part of the polymer in the blend forms continuous phase and the high viscous polymer confines to small domains.55 In the present case, it is apparent that in all compositions the high viscous PEI forms the continuous matrix containing low viscous PVH (for viscosity results see Figure 9A). This could be due to the high difference in the glass
transition temperature (Tg) between the two polymers (218 °C for PEI as against −33 °C for PVH), as measured using DSC, and the related details are provided under section 3.8. Upon addition of even 5% PVH to PEI (sample PHPI95), the small domain size of PVH is evident as tiny holes, and with further increase in PVH content to 10% (sample PHPI90), the holes appear stretched and the interfacial area becomes maximum (Figure 5B). The β-phase of PVDF upon interaction is known to exist in fibrillar form.56 Thus, it becomes easy to realize that during interaction, the PEI part of the matrix also got stretched. However, when the PVH content becomes 40%, the stretching is less noticeable due to insignificant interaction and the holes representing PVH still remains identifiable (Figure 5D). When the PVH content reaches to 60% (Figure 5E) and then to 80% (Figure 5F), the fibrillar form is more evident, since the PEI becomes minor phase (PVH rich phase). Thus, the SEM photomicrographs unambiguously proves that 12288
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content) until the PEI content is ≤80% and thereafter remains almost constant. The large positive change in I3 with PVH content shows that the o-Ps yield is strongly dependent on the PVH content, which appears to promote o-Ps formation. Since the −CF3 mainly contributes to the free volume,8,13,22 the o-Ps localized in a free volume hole will always find a fluorine atom in whose neighborhood it will annihilate.38,61 From what is known earlier, a marked reduction in I3 is expected owing to the stronger inhibiting effects caused by the presence of a large number of high electron affinity groups, which reduces the electron availability for Ps formation.38,61 Herein, with the increase in the PVH content in the blend (increase in fluorine atoms), the o-Ps yield also increases, thus showing “antiinhibition effect”, which is interesting. In the past, the addition of hexafluorobenzene (C6F6) to organic compounds was reported to show such an “anti-inhibition effect”.63 This has been explained as due to nondissociative electron attachment by C6F6 that allows more Ps to reach thermal energies and thus results to increased Ps formation.63 The work on poly(chlorotrifluoro ethylene) with three different vinylidene fluoride content also gives some indication toward “antiinhibition effect” in fluoro polymers.64 Here, we emphasize the recent reports from Sato et.al.41,42 that the exact nature of Ps interaction when localized in free volume holes that contain fluorine atoms on their cavity wall is still to be explored. The o-Ps intensity that depends on the probability of o-Ps formation is often considered by many authors that it may be proportional to the concentration of free volume holes.4,65 In the case of polymers having polar groups, it has been argued that Ps formation gets inhibited due to the strong electron withdrawing nature of the polar groups and hence I3 cannot be related to free volume number.60 As seen from the present results, with the increase in fluorine content in the blend, the oPs intensity also increases revealing that the fluorine atoms do not significantly affect the I3 value in this blend. Therefore, we have used I3 value to estimate the free volume fraction (Fv) of this blend using the relation
the interaction between PVH and PEI is high by showing maximum interface when the PVH content in the blend is ≤20% and is in agreement with the discussions made so far. Now, we will come to the center point of the discussion in this blend that relates the free volume and dielectric constant. 3.5. Positron Results. 3.5.1. PALS Results. 3.5.1.1. Discrete Lifetime Analysis. The three lifetime components from PALS spectra are attributed to positron and Ps annihilation in the blend. The shortest lifetime τ1 with its intensity I1 is attributed to p-Ps and free positron annihilation, while the second lifetime τ2 with its intensity I2 is from trapping of free positrons in defects and negatively charged polar sites and subsequent annihilation.41,42 The longest lived lifetime component τ3 with its intensity I3 is due to pick-off annihilation of o-Ps in the free volume sites. In the present study, we focus on the results of oPs pick-off annihilation only as it is related to free volumes in polymers.34 The semiempirical equation of Tao-Eldrup57,58 is used to calculate the mean “free volume hole” radius from the measured o-Ps pick-off lifetime (τ3) as τ3 = 0.5[1 − (R /R o) + (1/2π ) sin2π (R /R o)]−1 ns
(1)
This equation assumes spherical “free volume holes” of radius R with Ro = R + ΔR, where ΔR is the electron layer thickness. Nakanishi et.al.59 derived a value of ΔR = 0.1656 nm for molecular solids with known hole sizes. The average “free volume hole” size is found using the relation Vf3 = (4/3)πR3
(2)
The I3 represents the o-Ps formation probability.5,60 The changes in o-Ps lifetime (τ3), free volume (Vf3), and o-Ps intensity (I3) as a function of PEI content in the blend are shown in Figure 6A,B. The PVH has τ3 = 2.3 ns and I3 = 9.0%; and that for PEI has τ3 = 1.7 ns and I3 = 2.7%. The large free volume size in PVH (ca. 130 Å3) arises mainly from the asymmetric −CF3 group, since it creates sterical hindrances for a dense molecular packing.8,22,38,61 The lifetime values of PVH and PEI are in good agreement with those reported.38,62 The o-Ps lifetime shows an initial constancy and attains a maximum when the PEI content is 60% and starts to decrease significantly at the later stages (PEI ≥ 80%). The increase in lifetime with PEI content is due to increase in free volume which is an indication of noncompatibility between the two polymers.4 However, when the PEI content is ≥80% (PVH in minor phase), the free volume decreases significantly and reaches to ca. 70 Å3, and for this reason τ3 also decreases. From the combined FTIR, XRD, and TGA measurements, it is clear that, when the PEI content in the blend is ≥80%, significant dipolar interaction occurs between PVH and PEI, and the SEM shows increased interfaces in this range of PEI content. Because of this interaction, the available free volume and thereby the mobility of the polymer segments gets reduced. In a study on TiO2−PMMA nanocomposite, hydrogen bonding interaction was reported to decrease the τ3 value, which is in consonance with our results.60 The o-Ps intensity (I3) in the blend as a function of PEI content is shown in Figure 6B. A polymer that provides high probability for recombining with free electrons would exhibit less o-Ps intensity.60 In the present case, PVH provides lower probability of recombining with free electrons, which can be understood from its higher o-Ps intensity (I3 = 9.0%) compared to that for PEI (I3 = 2.7%). The I3 decreases continuously with increasing PEI content (I3 increases with increase in PVH
FV = CVf3I3
(3)
where C is the structural constant that can be obtained from other experiments.4,34 To understand the relative change in free volume, the Fv here is simply expressed as Fvr = Vf3I3 which is termed as relative fractional free volume (Fvr)65 and is used to discuss the results as relative free volume content. The variation of Fvr with PEI content (see Figure 6B) reveals that the relative free volume fraction continuously decreases with the increase in PEI content even when the o-Ps intensity remains almost constant at higher PEI compositions (PEI wt % is ≥80%). This is due to significant decrease in the average free volume size (see Figure 6A). It is known from SEM results that the maximum interface arises in this blend when the PEI content is ≥80%. The almost constant value of I3 in these compositions suggests that there is no significant contribution from interfaces to the observed I3 values (see Figure 6B). In other words, there are less chances that o-Ps would form at the interfaces. The lower value of I3 when the PEI content is ≥80% also suggests that Ps formation is strongly inhibited when dipole interaction is maximum, since the field due to dipole interaction seems to be not promoting Ps formation. 3.5.1.2. Continuous Lifetime Analysis (Free Volume Distribution). The free volume distribution is due to the fact that they result from chain folding and molecular architecture 12289
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which are of different sizes. The distribution of o-Ps pick-off lifetime obtained using the MELT program are shown in Figure 7 as relative intensity vs lifetime data.5 The corresponding “free
is in agreement with the results discussed so far. A schematic representation that reveals how the PEI content affects the change in free volume of PHPI blends is shown in Figure 8. 3.5.2. Doppler Broadening (DB) Results. In PALS, the three lifetimes representing the lifetime of p-Ps, free positron, and oPs can be well resolved. Even though DB spectra cannot be resolved into three components, the S-parameter that describes the width of the annihilation line itself provides qualitative information about the electronic structure in polymers.41,42 Generally, the narrow distribution is attributed to p-Ps part since it self-annihilates with low momentum electrons and this increases the S-parameter. In comparison to p-Ps, the momentum distribution of o-Ps is broadened because, the pick-off annihilation is influenced by electrons bound to the surrounding molecules in the free volume cavities and this decreases the S-parameter.35,38,39,41,42 Theoretically, due to spin degeneracy, upon Ps formation, 1/4 will be of p-Ps and 3/4 will be of o-Ps.34 Since S-parameter in a polymer is a measure of pPs while I3 is a measure of o-Ps, S-parameter would follow the trend of o-Ps intensity rather than o-Ps lifetime.35,39,41,42 The S-parameter increases with increase in PEI content in the blend while the W-parameter decreases with PEI content (see Figure 9A). In fact, an opposite trend is observed for S and W parameters when plotted against o-Ps intensity (see Figure S2 in the Supporting Information). The decrease of Sparameter with I3 could be due to (i) suppression of Ps formation, a process in which reduced contribution of p-Ps selfannihilation broadens the 2γ annihilation peak, and (ii) positron annihilation with high momentum electrons during pick-off process, which yields additional Doppler-broadening of the annihilated radiation.38,39,41,42 Thus, the S-parameter depends on both Ps formation and the kinetic energies of the electrons (surrounding the free volume cavity wall) with which the positrons annihilate. The decrease of S-parameter with increase in o-Ps intensity reveals that the broadening is caused more by the chemical environment through the kinetic energy of the annihilating electrons rather than through Ps formed.38,41,42 On recollecting the PALS results, the o-Ps formation is higher in the amorphous part (HFP) of the spur, and the free volume holes have been thought to arise mainly due to the presence of bulky −CF3 groups inferring that the walls of the holes contain fluorine atoms. When Ps resides in these holes, it would encounter electrons (chemical environment) from the surrounding highly electro negative fluorine atoms. The higher kinetic energy of the p-electrons of fluorine (F:2s2 2p5) would lead to a broader
Figure 7. Distribution of o-Ps lifetime (τ3) and free volume hole size (Vf3) for representative blend compositions.
volume hole” sizes are evaluated using eq 1 and are also shown at the top line for comparison. Please note that the average hole sizes obtained from MELT are close to those obtained from PATFIT analysis (see Table S1 in the Supporting Information) as reported by others.5 Generally for an immiscible blend, the distribution is quite broad compared to pure polymers, while a monotonical free volume distribution is observed in the miscible blend.4,5 In the present case, although, both the pure polymers have a narrow free volume distribution, it becomes broad with small increase in wt % of PVH. However, the pattern stays within the distribution limits of their pure polymers until the PVH content in the blend is ≤20%. This suggests good interaction between the two polymers4,5 Thereafter, the distribution pattern becomes more broad. The correlation between discrete and continuous lifetime analysis can be very much appreciated from the distribution pattern for PHPI60. Here, the distribution becomes highly broad and shows a significant upshift in lifetime, which partly lies above the upper limit of pure polymers and the discrete analysis also yields a maximum value of τ3 at this composition (see Figures 6A and 7). This is clearly a signature of noncompatibility between the two polymers5 and
Figure 8. Schematic illustration of arrangement of chains in PVH/PEI blends (a) when the PEI content is 20% and at this composition, we know that the interaction of PEI and PVH is insignificant. This incompatibility results in more free volume in the blend which in turn increases the chain mobility (Figure 6A). Additional factors such as, low Tg of PVH (Tg = −33 °C) and nonsymmetrical positioning of fluorine atoms in PVH also adds up to the free volume in the blend and contributes to the orientation and relaxation of dipoles.6,16 Thus, the dielectric constant increases. The dielectric constant remains almost constant with increase in free volume size until the size reaches ca. 125 Å3, and thereafter it steeply increases (see Figure 11B). Infact, positive deviation of dielectric constant can more clearly be evidenced when plotted against relative free volume fraction (see Figure S3 in the Supporting Information). In the case of PEI, the room temperature measurement corresponds to hundreds of degrees below Tg (Tg = 218 °C) due to which the molecular motion is severely limited and the contribution to dielectric constant comes mainly from electronic polarization.16 In order to further strengthen these findings on dielectric constant and free volume, we have studied the thermal behavior of the dielectric constant and the loss factor (tan-δ) of this blend. For many electronic applications, dielectric materials with stable dielectric constant across a large temperature range are preferred.7,9 Furthermore, the nature of variation of loss factor with temperature for a blend helps to know the Tg of
of electro negative fluorine atoms decreases the ε′ value owing to less electronic polarizability.8 On the other hand, fluorine atoms depending on its location and available free volume in the matrix can also produce additional dipole moment by increasing the orientation component of polarization.26 In the case of polymer blends, the dielectric constant at low frequencies arises mainly due to the contribution from electronic polarization, interfacial polarization, and/or dipole orientation.6,8,20,29,32,55 Interfacial polarization is commonly observed in heterogeneous systems like polymer blends and composites having different dielectric permittivities.55 This arises due to accumulation of mobile charges at the interfaces constituted by unlike phases.55 The dielectric constant when plotted against PEI content shows a continuous decrease until the PEI content is ≤80% and thereafter the change becomes insignificant (see Figure 11A). The pure PEI has ε′ value of 3.4 while that for PVH is 12.6 and these values are in close agreement with that reported.9,10,67 We know that there is a good intermolecular interaction between PVH and PEI with maximum interface when the PVH content in the blend is ≤20%. The almost constant value of ε′ in these compositions reveals that interfacial polarization is not contributing significantly to dielectric constant in this blend. On the other hand, the dipole orientation polarization arises due to redistribution of charges when a group of atoms with a net permanent dipole moment reorient itself in space fast enough to keep in pace with the oscillations of an applied alternating electric field.6,32 The process of dipole polarization accompanies the movement of polymer chain segments and the movement requires considerable free volume in the matrix. However, when the two polymers are compatible, it is difficult 12292
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Figure 13. DSC plots of PVH/PEI blend during (A) second heating and (B) second cooling for various PEI contents. The plots are shifted vertically for clarity.
associated with local bond rotations along the polyimide backbone.31 The other high temperature relaxation peak at ca. 229 °C is due to relaxation associated with Tg of PEI (αa‑PEI).67 From the discussions made so far, it is clear that intermolecular interactions exist between the polymers when the PEI content in the blend is ≥80%. The change in tan-δ of the two pure polymers with temperature is as an indication for compatibility of the blend.67 If the two Tgs of the individual polymers shift toward each other, it infers that the polymers are partially compatible and the presence of two Tgs at their respective positions represents that it is incompatible.5 This method is applicable for polymer pairs with well separated glass transition temperatures. Since the relaxation associated with βtransition of PEI and the αc of PVH overlaps, to identify the change in Tg clearly, tan-δ plots are made separately in the temperature range −100 to −10 °C (Tg region of PVH) and 120 to 250 °C (Tg region of PEI) for the blends with PEI content ≥80% along with the pure polymers (see Figure S4(A) and (B) in the Supporting Information). The tan-δ peak at 229 °C shows a slight downshift in temperature when the PVH content in the blend is ≤20% indicating sign of compatibility between the two polymers. However, in this composition range, the change in tan-δ peak of PVH could not be clearly ascertained due to low PVH content (see Figure S4 (A)). 3.8. DSC Results. The Tg is commonly deduced from DSC measurements4,5 and has been employed here to confirm the blend miscibility. The DSC traces for the entire blend composition during second heating and cooling are shown in Figure 13A,B, respectively, and the data are tabulated (see Table S2 in the Supporting Information). The PVH exhibits a weak transition at ca. −30 °C (marked as arrow in Figure 13A) due to its Tg and an endotherm at ca. 140 °C due to melting of PVDF crystals.11,30,51 The low value of Tg in PVH is due to the presence of −CF3 group that causes sterical hindrances to dense molecular packing and thus leads to high segmental mobility.38 The pure PEI shows a Tg at ca. 218 °C.54 The blends showed no evidence of Tg related to the PVH part while the Tg of PEI which is clearly evident in all the blend compositions, showed no significant change (see Table S2). The melting endotherm (Tm) showed a slight downshift when
pure polymers which would in turn help us to understand the blend compatibility.67 The plot of dielectric constant against temperature is shown in Figure 12A. The magnitude of ε′ depends on the quantitative alignment of the dipoles to the applied electric field. At room temperature, PVH part of the blend is in the rubbery phase (Tg = −33 °C) and PEI part is in the glassy phase (Tg = 218 °C). Thus, at very low temperatures (T < −35 °C), both the neat polymers of the blend are in the glassy phase and hence the molecular motions are frozen in place due to which the dipoles cannot move to align themselves to the applied electric field. When T > −35 °C, the chain mobility of PVH increases owing to availability of more free volume, and polar groups start to move in response to the applied electric field, which increases the orientation of the molecules and the dielectric constant.6,67 It is interesting to note that the ε′ value exhibited only a small change (3.4 to 3.94) when the PVH content in the blend is ≤20% in the temperature range of −100 to +100 °C (Figure 12A) again signifying good interaction between PVH and PEI. However, for the blends with PVH content >20%, the dielectric constant increases significantly, and at 120 °C, it reaches to nearly ten times its initial value. The variation of tan-δ (loss factor) with temperature for PVH resulted to two relaxation peaks: the first one at ca. −21 °C and the second one at ca. 22 °C. In an earlier dielectric study on PVDF measured at 1 kHz, two relaxation peaks were observed. The first one is centered at ca. −20 °C, and the second one is relatively broad, and its distribution starts from 20 °C and extends up to 120 °C, with a maximum at ca. 85 °C.68 The first relaxation peak (αa) was attributed to segmental motions in the amorphous phase associated with its Tg. The second relaxation (αc) is associated with the crystalline part and is due to imperfections developed with the chain during crystallization.68 Furthermore, if the crystals are completely of β-phase, then the αc relaxation does not exist.29 On extending these arguments to the present case for PVH, the first tan-δ peak seen at −21 °C is attributed to relaxation associated with Tg of PVH (αa‑PVH) and the second one seen at 22 °C is due to relaxations in the crystalline part (αc). The pure PEI exhibits two relaxation peaks one at 19 °C due to β-transition that is 12293
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the PEI content is ≥80% which again supports the interaction between the two polymers. In an earlier study on the blend of PEI with semicrystalline poly(butylene terephthalate) (PBT), the authors have observed depression in the Tm of PBT and was attributed to good miscibility between the two polymers,69 which is in consonance with the present results. The cooling curve for PVH shows an exotherm at ca. 103 °C due to crystallization of PVDF and represents the crystallization temperature (Tc). This temperature remains almost unchanged until the PEI content is 60% (sample PHPI60). Interestingly, when the PEI content is ≥80%, the Tc first shows a downshift and another small exotherm starts to appear at ca. 58 °C having enthalpy of crystallization (ΔHc) in the range of 0.58−1.8 J/g (see Figure 13 B and Table S2 in the Supporting Information). We are yet to understand the link between this delayed crystallization and the dipole interaction observed in this blend. 3.9. DMTA Results. It is known that fluorinated polyimides suffer from low mechanical strength.23 In order to understand the change in stiffness of the blends, DMTA measurements were done. In a DMTA measurement, the storage modulus (E′) is a measure of stiffness of the material and is plotted as a function of temperature for various compositions of the blend (see Figure 14).
Thus, both PVH and PEI remain in the glassy phase in the first temperature zone, while in the second zone, the PVH attains a rubbery phase but PEI continue to remain in the glassy phase. The results revealed that, even though PVH has poor storage modulus at room temperature, the blend attains stiffness on par with PEI, at higher PEI content. This increased stiffness is particularly important when PVH is used as a polymer electrolyte.9,12 When PVH is used in such applications, the liquid electrolyte uptake decreases the mechanical strength of PVH. Wang et.al. have reported the use of a microporous polyolefin to improve the mechanical properties of PVH.70 Herein, we propose blending of PVH with PEI to be another option to improve the mechanical property of PVH. In summary, when the PEI content is ≥80% or PVH content is ≤20%, the blend acquires improved thermal stability and stiffness compared to PVH. On the other hand, it obtains high thermo-oxidative stability and less viscosity compared to PEI. Thus, we propose PVH/PEI (20/80) to be the optimum composition for use of this blend. It is known that introduction of free volume to polyimides and related materials results in a decrease in the dielectric constant.6,17,18,31−33 The free volume in this blend shows a positive variation with dielectric constant, which is interesting and marks this blend as quite different in comparison to other fluorinated polyimides and their blends.
4. CONCLUSIONS In this work, we report the structure−property relationship regarding the effect of free volume, fluorine content, and dielectric constant in a polyimide blend. The study reveals that a dipolar type interaction exists between PVH and PEI and has maximum interface in their blend when the PVH content is ≤20%. The origin of the dielectric constant in this blend is identified to be mainly from dipole orientation polarization owing to the availability of more free volume. The positive variation of dielectric constant with free volume content is distinct from the generally observed behavior in polyimides and their blends studied so far. The present results would help with a more effective design of polymeric electrical charge storage devices. Furthermore, the results from positron measurements show that increase of o-Ps fraction with fluorine content, leading to “anti-inhibition effect”, has been observed for the first time in a polymer blend.
Figure 14. Plot of storage modulus (E′) of PVH/PEI blend with various PEI content.
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ASSOCIATED CONTENT
S Supporting Information *
From the nature of variation of E′ with temperature, it can clearly be separated into two temperature zones. First one is −100 °C ≤ T ≤ −50 °C and the second is −51 °C ≤ T ≤ 120 °C. In the first zone, the storage modulus of PVH is high (ca. 4.0 GPa) and that of PEI is relatively low (2.8 GPa). The E′ value of the blends for various compositions lies between that of PVH and PEI. Upon heating PVH, the modulus drops drastically at ca. −50 °C as it passes its Tg and goes to rubbery phase. Upon heating the blends, with the increase in PEI content, the drop in modulus seen around this temperature becomes low. Thus, in the second zone, an opposite behavior is observed. Here, the E′ value of PVH remains the least, and with increase in PEI content, the E′ value increases. When the PEI content is ≥80%, the modulus remains almost stable even beyond 120 °C and thus permit measurements above 120 °C. When the temperature is close to 220 °C, the E′ drops again, this time steeply, signifying the Tg of PEI. Please note that for blends with PEI content
