Influence of Microstructure on the Nanomechanical Properties of

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B: Fluid Interfaces, Colloids, Polymers, Soft Matter, Surfactants, and Glassy Materials

Influence of Microstructure on Nanomechanical Properties of Polymorphic Phases of Poly(vinylidene Fluoride) Suresh Guduru, Sanjay Jatav, G Mallikarjunachari, M. S. Ramachandra Rao, Pijush Ghosh, and Dillip Kumar Satapathy J. Phys. Chem. B, Just Accepted Manuscript • DOI: 10.1021/acs.jpcb.8b05972 • Publication Date (Web): 20 Aug 2018 Downloaded from http://pubs.acs.org on August 20, 2018

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The Journal of Physical Chemistry

Influence of Microstructure on Nanomechanical Properties of Polymorphic Phases of Poly(vinylidene fluoride) G. Suresh,†,‡ Sanjay Jatav,†,‡ G. Mallikarjunachari,¶ M. S. Ramachandra Rao,§ Pijush Ghosh,¶,‡ and Dillip K. Satapathy∗,†,‡ †Soft Materials Laboratory, Department of Physics, Indian Institute of Technology Madras, Chennai 600036, India. ‡Center for Soft and Biological Matter, Indian Institute of Technology Madras, Chennai 600036, India. ¶Department of Applied Mechanics, Indian Institute of Technology Madras, Chennai 600036, India. §Nano Functional Materials Technology Centre and Material Science Research Centre, Indian Institute of Technology Madras, Chennai 600036, India. E-mail: [email protected] Abstract Poly(vinylidine fluoride) (PVDF) is a semi-crystalline polymer which is known to exist in several polymorphic phases, namely, α, β and γ. Each one of these polymorphic phase is characterized by unique features such as spherulite formation in the case of α- and γ-phases, and the presence of large piezoelectric and ferroelectric activity in β-phase. Despite being widely used as thin coatings in sensors, lack of reports on nanomechanical properties suggests that investigation of mechanical properties of

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PVDF, let alone those of its polymorphic phases seems to have evaded the sight of research community. Herein, we report the nanomechanical properties of α, β and γ phases of PVDF. The modulus and hardness values were evaluated from nanoindentation experiments; it was found that the electroactive β-phase is the softest among the three polymorphic phases. This result was further confirmed by scratch experiments. We have attempted to establish correlation between the microstructure and nanomechanical properties of these phases. This work sheds light on the mechanisms responsible for the observed mechanical behavior, and the role of tie molecules and amorphous content in providing flexibility to the polymer.

Introduction PVDF has captured the interest of researchers not only because of its sizable piezoelectric, pyroelectric and ferroelectric properties but also due to its coexistence in several polymorphic phases namely α, β, γ and δ. 1,2 The chain conformations of these phases are shown in figure 1. Thermodynamically, α-phase is the most stable polymorph of PVDF at ambient conditions, which is obtained from melt crystallization of PVDF. 3,4 Chain arrangement in β-phase unit cell results in a non-vanishing dipole moment, 5,6 rendering β-phase crystal ferroelectric. The polymer chains in γ-phase have a conformation that lies between that of the α- and βphases. PVDF is one promising functional polymer extensively used in applications ranging from energy transductions 7,8 to filtration membranes 9,10 and non-volatile FE-RAMs 1,11 to space mirrors. 12 PVDF is widely studied for its outstanding electroactive properties, chemical inertness and high dielectric constant compared with other polymers. 13,14 The physical properties of PVDF and content of different polymorphic phases can be controlled by varying synthesis conditions. 15,16 From the above discussion, it becomes apparent that electrical properties and polymorphism in PVDF have been studied quiet extensively. Mechanical properties of materials play a pivotal role in determining the function and realizing smooth operation of devices. For example, in piezoelectric materials, variation in mechanical prop2


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erties is immensely significant because the functionality of the material is strongly hinged to it. In the realm of biology, it is being touted that piezoelectricity can also be viewed as an extended property of animate matter like living tissues and bones, that plays a vital role in numerous physiological phenomenon. 17 Very recently, it was found that the elasticity of nanoparticles can strongly influence the drug uptake by certain kind of tumor cells. In their study on human breast cancer cells, Guo et. al. found that the intake of mechanically softest nanoparticles was 80% greater than their stiffest counterparts. 18 Variation in mechanical properties is very also crucial for tissue engineering applications. For example, the growth of bones and tissues is markedly affected by the stiffness of the scaffold which can be simulated mechanically. 19 Since, PVDF is used in applications like coatings and sensors, where the material is in contact with other surfaces, it is puzzling to see the lack of literature on mechanical properties of PVDF.

Figure 1: Chain conformation in Polymorphic phases of PVDF in (a) α-phase (b) β-phase and (c) γ-phase. α-phase is non-polar, whereas both β- and γ-phase have a non-zero dipole moment of the unit cell. The mechanical properties of materials are strongly correlated with the underlying microstructure. 20,21 For example, in our previous study we found that the hardness of PVDFP(VDF-TrFE) blends is intimately linked to their morphology. The hardness values show a marked decline when the blend morphology changes into fiber-shaped crystals. 16 Therefore, understanding crystallization in polymers becomes imperative for tailoring their mechanical properties. We have examined the nanomechanical properties of polymorphic phases of PVDF by employing nanoindentation (NI) technique. One of the merits of this technique is that it

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allows the application of very low loads (ranging from nN to µN) to the sample and measures the small displacements. This technique is very powerful in the sense that it allows one to obtain information on localized material properties and map the mechanical properties with a very high spatial resolution. 22 NI scores over bulk mechanical characterization techniques since, hardness evaluated from NI is much more sensitive to the surface features of the materials. With NI, mechanical properties of the surface 23 and the effect of substrate on the mechanical properties of the material can be probed. This technique also provides the means to study failure in a material at a molecular level. 24 Since, PVDF finds application in (thin) coatings, NI is the most appropriate technique to study its (nano)mechanical properties. Recently, we reported on the enhancement of mechanical properties of drop-casted PVDFP(VDF-TrFE) blend films by increasing the crystallization temperature. 16 Oliveira et al. performed nano and micro indentation tests on PVDF and found that in low-load regimes (∼100 to 5000 µN/s load rates) the modulus values decrease with loading rates, while at higher loads the modulus values tend to increase with increasing loading rates. 25 According to the authors, during nano-scale loading, small volume deformations and local heating effects lead to a decrease in viscosity along with a subsequent drop in elastic modulus. But they did not explore the effects of much lower (< 100 µN/s) load rates on the mechanical behavior of PVDF. While these aforementioned studies shed light on the influence of crystallization and testing conditions on the over all mechanical properties of PVDF, there are no reports which discuss the mechanical properties of α, β and γ phases of PVDF exclusively. To the best of our knowledge there is a meager body of literature which focuses on the structure-mechanical property correlation of individual phases of PVDF and their respective morphologies. The objective of this work is to compare the nanomechanical response of α, β and γ polymorphic phases of PVDF and determine the influence of spherulitic microstructure on the same at nanometer length scales. This article is organized in following sections. The first part deals with synthesis methodology and experimental details. The following section discusses the identification and quan-

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tification of different phases of PVDF and their morphology. Next section describes results of nanoindentation experiments, their correlation with the observed morphology and finally closes with conclusions.

Experimental Section PVDF in powder form with average molecular weight MW ∼534,000 g/mol and N,N Dimethylformamide (DMF) of ACS reagent grade were procured from Sigma-Aldrich India. These chemicals were used as received. One gram of PVDF powder was added to 5 ml of DMF solution and subjected to ultrasonication for homogeneous dispersion and mechanically stirred till the complete dissolution of the polymer. Glass substrates were first washed with soap solution in DI water, followed by cleaning with iso-propyl alcohol and acetone. Then, glass slides were further rinsed several times with DI water to remove any traces of isopropyl alcohol and acetone and also flushed with dry nitrogen gas for complete removal of water from surface. PVDF films were obtained by spin-coating the polymer solution on to the clean glass substrates. Nearly 300 µl of polymer solution was spun at 400 revolutions per minute. Solvent evaporation was performed at fixed temperatures i.e. at 210 ◦ C for 10 minutes for obtaining α phase rich PVDF. 26 This sample is referred as α-210 in rest of the manuscript. Similarly, in order to obtain β phase rich PVDF films, solvent evaporation was carried out at 60 ◦ C for 10 hours(referred as β-60). 27 All the samples were allowed to cool naturally. Thickness of these films were measured to be about 4 µm. X-ray diffraction measurements were carried out on Rigaku Smartlab X-ray diffractometer using CuKα radiation at room temperature. Fourier transform infrared spectroscopy (FTIR) spectrum of the films were recorded on a Bruker Alpha-Platinum ATR spectrophotometer in attenuated total reflection (ATR) mode with a resolution of 4 cm−1 with each spectrum averaged over 32 scans. Dielectric relaxation spectroscopy (DRS) measurements

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were performed on Novocontrol Alpha-N High Resolution Dielectric analyzer with active sample cell ZGS as test interface. These measurements were carried in the frequency range from 1 to 107 Hz at room temperature. Surface morphology of these polymer films was investigated using electron microscopy on FEI-Quanta 200F high resolution scanning electron microscope (HRSEM). Polarized optical microscopy (POM) was performed using Axioskop 2 MAT, Cal Zeiss. The nanoindentation experiments were carried out using the Berkovich indenter (Hysitron TI Premier) tip made of diamond (Young’s modulus, E= 1140 GPa, Poisson’s ratio, ν = 0.07). Appropriate load function (loading 20 sec, holding 10 sec and unloading 10 sec) was used for all the samples.

Results Identification of polymorphic phases of PVDF Due to the inherent differences in the arrangement of atoms and polymer chains in α- and β-phases, these phases can be identified based on their respective crystal structures and conformations which can be probed by using XRD and FTIR measurements, respectively.

Crystal structure Since the crystal structures of α- and β-phases are different, 5 X-ray diffraction can be used to identify and differentiate these polymorphs. X-ray diffractograms recorded on α-210 and β-60 PVDF samples are shown in figure 2 (a). The reflection peaks at 17.7◦ , 19.9◦ and 26.8◦ are seen due to diffraction from (100), (110), and (021) set of planes, respectively (present in black profile), which corresponds to the α-phase of PVDF. 4 The peak at 20.2◦ corresponds to the sum of diffractions from planes (110) and (200) (present in red profile) which is the characteristic peak for electroactive β-phase. 4,6 In addition, an amorphous hollow can also be observed at slightly lower angle for β-phase (present in red profile). From these observations it can be established that each film predominantly contains an individual phase of PVDF; 6


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nevertheless they can still have small fractions of other phases.

Figure 2: (a) X-ray diffractograms of polymorphic phases of PVDF (b) FTIR spectrum of polymorphic phases of PVDF. Both X-ray diffractrograms and FTIR spectra are shifted vertically for visual clarity. Vertical dotted lines at specific values are guide for eye to observe evolution of the peaks. From x-ray diffractogram it is clearly seen that crystallinity of PVDF improves with increasing solvent evaporation temperature.

Difference in conformations FTIR is a complementary technique for quantifying and distinguishing between the different phases present in PVDF. There exist different absorbance bands which are peculiar to these individual phases. 4,6 PVDF crystallized from temperatures above melting point preferentially crystallizes in non-electroactive α-phase. 3 From the FTIR spectrum shown in figure 2 (b), the absorbance bands at 615, 762, 794, 975, 1144, 1210 and 1383 cm−1 corresponds to α-phase. 6,28 Absorbance bands at 840 and 1400 cm−1 indicate the presence of electroactive β-phase. 6,29,30 The amount of different phases present in films prepared at 210 and 60 ◦ C were calculated using the procedure described by Martin et al. 6 The amount of α- and β-phase in 210 ◦ C and 60 ◦ C treated samples were found to be nearly 75% and 83% respectively. The remaining fractions in each film can be associated with mixture of other phases and amorphous phase. 31 Hence, FTIR analysis confirms that PVDF films with predominantly α-

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and β polymorphic phases have been synthesized.

Dielectric properties of polymorphic phases of PVDF Room temperature dielectric relaxation spectrum (DRS) measurements were carried out in order to investigate the dielectric properties of individual phases of PVDF. The variation of real part of dielectric constant, 0 , with frequency is shown in figure 3.

Figure 3: Dielectric relaxation spectrum of polymorphic phases of PVDF with varying frequency real part of modulus. Electroactive β-phase shows highest value of due to its non-vanishing dipole moment. From the figure 3 it can be seen that dispersion of 0 shows a similar trend for both αand β-phases of PVDF. β-phase being the most polar among all the phases of PVDF shows the highest value of 0 followed by γ-phase which also has non-zero resultant dipole moment (not shown in the figure). Similar results were obtained by Gregorio et.al.,. 15 0 decreases with increasing frequency as it becomes increasingly difficult for the dipole to follow rapidly varying electric field. Thus, the contribution of dipolar groups towards the permittivity reduces continuously as the frequency increases. High value of 0 observed at low frequencies are known to occur due to ionic dc conductivity contribution, which results in interfacial or spatial charge polarizations. 15 As the frequency goes past 103 Hz, reduction in value 8


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of 0 occurs due to the dipolar relaxation of CF2 - CH2 molecules of the PVDF polymer chain whose characteristic relaxation frequency range lies from 101 to 106 Hz. 32 From DRS measurements, we can clearly see the difference in dielectric behavior of the two phases. This result further reinforces the fact that α- and β-phases have been synthesized.

Microstructure As mentioned earlier, microstructure of a material strongly influences their mechanical properties. Therefore, the morphology of polymorphic phases of PVDF was closely examined using polarized optical microscopy (POM) and scanning electron microscopy (SEM). It is well established that when semi-crystalline polymers are grown from melt, the repeated branching and splaying of lamellae results in the formation of ball-like spherical structures called spherulites. 33,34 Spherulites being birefringent are sensitive to polarized light and hence, POM becomes the instrument of choice for investigating spherulite formation in different phases of PVDF.

Figure 4: Polarized light microscopy images of (a) α-phase (b) β-phase. Spherulites can be clearly seen in the case of α-phase. Generally, melt-crystallized semicrystalline polymers form these spherulitic patterns. POM images of α- and β-PVDF can be seen from figure 4. Spherulites can be clearly observed in case of α-PVDF (fig. 4(a)) whereas, for β-PVDF, the formation of spherulites is not so evident. The size of spherulites in α-PVDF was found to be nearly 40 µm in size. 9



Figure 5: SEM micrographs showing surface morphology of (a) α- and (b) β-phases of PVDF. Variation in size of the spherulites (surface morphology) as described in morphology section can be distinctly seen for different polymorphic phases. A closer examination of surface morphology of these samples was carried out using scanning electron microscope. The SEM micrographs of these phases are shown in figure 5. For α-phase, large domains nearly 40 µm in size are observed (figure 5(a)). From the inset of figure 5(a), it can be seen that dendrites (lamellae) extend radially outward from nucleation points. Micron to sub-micron sized grains were observed in β-phase along with the presence of few pores which can be seen clearly from figure 5 (b). Similar morphology for β-phase was observed in our previous work where PVDF film was prepared at 70 ◦ C. 35 At low temperatures due to slow rate of evaporation of solvent the formation of pores becomes favorable because of vapor-induced phase separation which leads to the formation of incomplete crystallites. 36 Atomic force microscopy (AFM) further reveals the surface topography of different polymorphs of PVDF at high spatial resolution. The AFM images further confirm that α-phase has a very smooth surface whereas β-PVDF exhibits a more coarse and rough morphology. The topography of β-PVDF comprises of many undulations which appear like hills and valleys. In the case of polymers, structure formation is governed by kinetic aspects rather than by equilibrium thermodynamics. 37 As a result, the structure that develops at a given temperature is the one with the maximum growth rate instead of that with the lowest free energy. 38 Therefore, being kinetically controlled, structures of semi-crystalline polymers are 10


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always affected by the processing conditions like cooling rate. 39 In our case, the difference in processing condition stems from the different solvent evaporation temperatures chosen for obtaining α and β phases, respectively. As can be seen from the above POM images (figure 4) and SEM micrographs (figure 5), the morphologies differ significantly due to difference in solvent evaporation temperatures. The reason for this lies in the fact that evaporation rate of solvent has a significant bearing on the final morphology of PVDF. 36 Difference in grain sizes of α- and β-phases is probably due to the decrease in the nucleation rate and increase in the growth rate of the spherulites with increasing temperature. Our observations further lend support to the growing body of evidence that processing conditions alter the morphology of semi-crystalline polymers drastically. 3,40,41

Mechanical properties of α−phase and β−phase Since, PVDF is widely used in coatings, membranes, sensors and actuator applications, characterization and knowledge of mechanical properties becomes a crucial step in ensuring seamless operation of the device 42 and prevention of unwanted failures. For examining the difference in response of α- and β-phases of PVDF to external force, nanoindentation (NI) and nanoscratch (NS) experiments were conducted.

Quasistatic Nanoindentation Initially, different loads ranging from 500 µN to 2000 µN were applied on all the samples to determine the bulk properties. Care was taken to minimize the substrate effect. Ten separate indents were performed at each load. Figure 6 shows the dependence of reduced modulus, Er on applied load for both α- and β-PVDF. The figure indicates that at a load of 2000 µN, Er attains a steady value, free from any surface effects and noises. Therefore, for all subsequent measurements and comparisons, 2000 µN was applied. The hardness, H and Er values were evaluated from the load-displacement plots applying Oliver-Pharr equations, 43

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H=

Pmax Ac

+ ... + C4 h1/4 Ac = f (hc ) = C1 h2c + C2 hc + C3 h1/2 c c √ π S √ , Er = 2β Ac

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(1)

(2) (3)

where, H is the hardness at nanometer length scale; Pmax is the maximum load; C1 , C2 , C3 , C4 and β are constants and Er is the reduced Young’s modulus; Ac is the contact area; hc is the contact height and S is the contact stiffness of the material as expressed in above equations. The selection of appropriate load rate plays an important role while determining the mechanical properties of visco-elastic materials. Series of experiments done with different load rates, however, indicate that the mechanical properties of PVDF are not very sensitive to load rate (see fig. 7). Hence, load rate of 100 µN/s was chosen for further experiments.

Figure 6: Modulus value dependence with varying loads for (a) α- and (b) β-phases. The sizes of the residual impression of contact during indentation with respect to the microstructural features of α- and β-PVDF are shown in figure 8 and the corresponding representative load-displacement curves are shown in figure 9. Modulus is calculated from 12


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Figure 7: Dependence of modulus on rate of loading for (a) α-phase and (b) β-phase. the slope, S of initial unloading segment of the load-displacement curve, along with equation (2) and (3). Table 1 indicates the average value of Er , and H obtained from at least 10 indents, from different locations, uniformly distributed over the sample. Interestingly, the modulus value for these phases do not vary so much, however by comparing the nanoscale hardness values of these phases it can be seen that α-phase appears to be hardest, indicating that this phase is more resistant to plastic deformation. β-phase appears to have lowest resistance to plastic deformation as indicated by the hardness values. Figures 9 and 10 also indicate that the maximum depth attained by β-phase is larger compared to other phases. Table 1: Modulus and hardness of individual phases of PVDF. Material α β

Er (GPa) H (MPa) 3.4 ±0.1 200 ±5 3.7 ±0.5 150 ±25

The horizontal segment (inset of fig. 9) of a load-displacement plot is a good indicator of material creep at nanometer length. At a sustained load, the β-PVDF is observed to undergo larger deformation compared to α-phase. This difference in response to the applied load primarily originates from different molecular structures of α- and β-phases of PVDF as indicated in figure 1. From a material design perspective it is essential to understand 13



Figure 8: Scanning probe microscopy images showing surface morphology after NI experiments on (a) α- and (b) β- phases of PVDF. Clear impressions can be seen in the case of α phase whereas for β-phase the indenter marks are rather diffuse and not so clearly visible. Size of the scan area is 75µm x 75µm.

Figure 9: Representative load-displacement plots of polymorphic phases of PVDF. Inset shows area corresponding to elastic and plastic energies. The horizontal segment of the load-displacement plots show material creep, which is greater for β-phase. the amount of incident energy a material can dissipate in the form of plastic deformation before failure. The area under the load-displacement plot (inset in figure 9) can be used to determine the plastic energy (EP ) and its elastic counterpart (EE ). From table 2, it can be seen that largest amount of energy (EP + EE = E) is expended in indenting the β-phase. As indicated by the ratio of EP /E, it appears that the microstructure of β-PVDF comprising of very small sized domains (grains) has more ability to dissipate energy through irreversible

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deformation than the α-phase. Surprisingly, the order of plastic energy stored in these phases resembles the order of dipole moment values of these phases, where β-phase has the largest dipole followed by γ-phase (not shown here) and α-phase is non-polar. Table 2: Mechanical energy components of individual phases of PVDF measured at 2000 µN peak load and 100 µN/s load rate. Phase α β

Elastic Energy (EE ) (10−10 J) 4.5 4.3

Plastic Energy (EP ) (10−10 J) 0.7 2.1

Plasticity index (Ψ) (EP /EP + EE ) 0.13 0.33

Interestingly, the standard deviation in Er and H for β-phase amounts to 13.5% and 16.7% respectively, which is quiet high when compared to α-phase (∆Er = 3%; ∆H= 2.5%) as can be inferred from table 1. In order to further probe the role of microstructure in determining the mechanical response, nearly 100 indents were performed with a maximum load of 2000 µN and load rate of 100 µN/s on both α and β-phase phases of PVDF. These loaddisplacement plots are shown in figure 10 along with the distribution of maximum depths. The width of the band of load-displacement plots obtained from 100 indents represents the deviations observed in Er and H earlier. The width of this band is least for α-PVDF; greatest for β-PVDF. These deviations in Er and H are captured quantitatively in the distribution of maximum depths plots shown in the bottom panels of figure 10. The distribution is found to be very narrow for α-phase, lying mostly between 650 to 750 nm, whereas for β-phase the spread lies between 550 to 1050 nm. The observed spread in maximum depths could be due to two probable reasons. Firstly, due to the variation in morphology of different phases of PVDF. The morphology of β-phase is quite distinct from α-phase. β-phase mostly comprises of small sized grains with very uneven topography. Therefore, while performing NI tests the indenter might not be able to make uniform contact with the surface leading to non-uniform distribution of applied load. The other reason might be the presence of fractions of other phases i.e. α- and γ-phase components in the β-phase, because of this, there could be deviations in the penetration 15



Figure 10: Load-displacement plots from 100 indents NI experiment of PVDF (a) α-phase and (b) β-phase. Their corresponding histograms are shown in the adjacent column. β-phase exhibits the widest distribution of maximum depths. The least spread is seen in the case of α-phase. depths. The morphology also explains why deviation in mechanical properties of α-phase is the least or why the distribution of maximum depth is narrow. Due to its pristine morphology comprising primarily of large sized grains (spherulites) with even topography, the indenter can impinge perfectly on the surface of α-PVDF. This translates to uniform application of load on the polymer which is also reflected as a clear residual contact impression on the surface of α-PVDF (figure 8(a)). The effects of this morphology and the reduced (comparatively) ability to resist plastic 16


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Figure 11: Representative load-displacement plots of β-phase. The sudden changes in the slopes of loading curves are shown as dotted circles. This could be due the impinging of indenter tip on structural irregularities like grain boundaries. deformation are also reflected through the pop-ins (circled) observed in the elasto-plastic loading phase of the β-phase as indicated in figure 11.

Scratch response Unlike nanoindentation experiments, during nanoscratch experiments, a lateral load is also applied along with the normal load. The scratch, several microns in length, thus, allows measurement of averaged mechanical response of the material. With scratch technique, it is possible to probe the inhomogeneities in mechanical response of the material, both laterally, and across the thickness. Generally, small indentation depths assess superficial layers of the polymer surface and features that are of the same scale as the size of the indenter. Therefore, for obtaining the response of the material averaged over a larger area of the surface, scratch experiments were performed on these polymorphic phases of PVDF. Nanoscratch experiments essentially probe the friction offered against the indenter tip while scratching. During scratch experiments, an optimized load of 2000 µN was applied and a 20 µm long scratch was etched on the surface of the polymer film. The plots of lateral displacement against normal displacement of the indenter tip for the both α and β polymorphic phases

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of PVDF are shown in figure 12. The friction coefficients were found to vary in a manner similar to figure 12 for α- and β-phases, respectively.

Figure 12: Normal displacement against lateral displacement plot of the indenter tip obtained from scratch measurement performed on α and β phases of PVDF. SPM images of these phases after scratch are also shown in the insets where the scan size of each of image is 50µm x 50µm. From figure 12, it can be seen that the depth (normal displacement of indenter tip) attained with the progress of the scratch is different for α- and β-phases of PVDF. For α-PVDF an almost consistent depth of 750 nm is maintained throughout the length of the scratch, whereas for β-PVDF the depth increases continuously in an oscillatory manner throughout the length of the scratch. The fluctuations observed in depth profiles of β-PVDF are due to the presence of surface inhomogeneities in the form of grain boundaries. Whenever the indenter tip passes through a grain boundary, it falls into the depression region present between the adjacent grains. This depression region can be clearly seen from the SPM images shown in the inset of figure 12. Therefore, in β-PVDF, due to the presence of very small sized grains and greater number of grain boundaries the depth profile shows oscillatory behavior as compared to α-PVDF. Another contributing factor can be the rough morphology of βphase as discussed previously. The continuous increase in depth along the scratch length 18


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observed in β-phase as opposed to α-PVDF means that the packing of polymer chains along the z-axis (normal to the surface of films) is markedly different from the phases prepared at high temperatures (α-210). From the scratch experiment it is again confirmed that despite it’s rough surface morphology, β-phase is softer than the other two phases.

Comments on γ-phase The authors’ have studied the mechanical response of γ-phase in detail and its properties were found to be intermediate to α- and β-phase. This is expected from its morphology. The results for γ-phase are not shown here for the purpose of attaining greater clarity in presenting the data.

Discussions The fundamental structure of semicrystalline polymers comprises of platelike crystallites with thicknesses in nanometer regime, embedded in liquid matrix. 38 When crystallized from melt, the continuous branching and splaying of the crystal lamellae (platelike, crystallites with curved edges) results in the formation of spherulites. 33,38 Considering the complexity of microstructure of semicrystalline polymers it becomes apparent that the structure-property relationships are also not so simple. 21 Hence, the simple use of chemical structure to predict mechanical properties, has severe limitations. The reason being, generally the structure of the individual molecules contributes only indirectly to collective mechanical properties, which arises from collective response of molecules. The supramolecular structure, i.e., the physical arrangement of the molecules with respect to one another being of more direct consequence. In effect, the mechanical properties are supramolecular structure sensitive in case of crystalline polymers; for example, a state of brittleness to one of toughness or ductility can be accomplished by suitable rearranging the polymer molecules within the same sample. 44

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We suspect that the observed difference in the mechanical properties of α- and β-phases arise from the arrangement of domains and amorphous regions in these phases; which essentially is function of the synthesis conditions. The mechanical properties of polymers are also influenced by the size of spherulitic domains. Just like nanocomposites, where particle distribution plays a key role in determining the properties, mechanical properties are also influenced by the spatial distribution of these spherulitic domains. Because of the inherent size difference between the domains of α- and β-phases (figure 5), they are very likely to show disparate shearing behaviors. This shearing behavior depends primarily on the surface to volume (S/V) ratio of these domains; which is much greater for β-phase. Another aspect of the structure of crystalline polymer is that the inter-spherulitic (domains) regions, and the inter- and intra-lamellar regions are filled with amorphous phases of polymers. The amorphous region also comprises of tie molecules; that provide molecular links between adjacent lamellae. 21,44 These tie molecules save semicrystalline polymers from being extremely friable. 44 On a molecular level they affect the mechanical properties by distributing the stresses more uniformly between all the lamellae; thus, conferring ductility to the polymer. 21 Due to the presence of large number of domains (of the order of 1 µm or even less) in β-phase, the amorphous regions are far more frequent between the grains as compared to that of α-phase. Therefore, it would be reasonable to assume the presence of large number of tie molecules between the adjacent domains (and lamellae) in β-phase. Hence, we conjecture that presence of large number of tie molecules renders to β-phase its soft nature. Since, β-PVDF consists of a larger fraction of surface covered by grain boundaries (due to high S/V ratio), the indenter tip (tip diameter of 150 nm) has a higher probability of impinging on domain boundaries as compared to α-phase (figure 13). The plastic deformation in these grains arise due to the chain moments both in crystalline and amorphous regions. 25 From the vantage point of free volume considerations, the softer nature can also be explained by assuming that greater free volume exists in amorphous regions, which implies that it can easily undergo plastic deformation resulting in relatively lower hardness value for β-phase.

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Figure 13: Schematic of arrangement of crystalline and amorphous content in the (a) αand (b) β-phases of PVDF with respect to the indenter tip dimensions (∼150 nm). In ’a’ indenter indents on the crystalline spherulite, whereas in ’b’ it indents on the inter-granular (amorphous) region.

Conclusions In this article, we have presented a detailed investigation of mechanical properties of α- and β-phases of PVDF . From NI experiments it was found that modulus values of these phases do not differ significantly from each other, whereas their hardness varies considerably; β-phase being the softest. The examination of large deviations in Er and H for β-phase reveal the role of surface morphology in determination of mechanical properties by NI. These deviations in Er and H values are smaller for bigger sized grains and vice versa. Bulk mechanical response of α- and β-phases evaluated by scratch experiments also establishes the fact that β-phase is softest. The difference in mechanical properties among these phases of PVDF primarily arises from their distinct and unique microstructure. The softer response of β-phase has been attributed to the large number of tie molecules present between the lamellae.

Acknowledgement DKS acknowledges the financial support from the Department of Science and Technology under DST-SERB, grant no.: EML/2016/003976

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Figure 14: TOC Graphic.

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Influence of Microstructure on the Nanomechanical Properties of

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