One-Step in Situ Modification of Halloysite Nanotubes: Augmentation

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One step in-situ modification of Halloysite nanotubes: Augmentation in polymer-filler interface adhesion in nanocomposites Purabi Bhagabati, T K Chaki, and Dipak Khastgir Ind. Eng. Chem. Res., Just Accepted Manuscript • DOI: 10.1021/acs.iecr.5b01043 • Publication Date (Web): 17 Jun 2015 Downloaded from http://pubs.acs.org on June 19, 2015

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Industrial & Engineering Chemistry Research

One step in-situ modification of Halloysite nanotubes: Augmentation in polymer-filler interface adhesion in nanocomposites Purabi Bhagabati, Tapan K. Chaki*, Dipak Khastgir Rubber Technology Centre, Indian Institute of Technology Kharagpur, Kharagpur 721302, West Bengal, India *Corresponding Author: Tapan K. Chaki (E - mail: [email protected])

ABSTRACT: Nanocomposites of chlorinated polyethylene (CPE)/ ethylene methacrylate copolymer (EMA) (60/40 ratio) with augmented polymer-filler interface adhesion were prepared by following a facile covalent modification of halloysite nanotubes (HNT). The covalent modification of the outer and inner lumen structures of HNT were pursued by two different methods. In-situ ring opening polymerization of ε-caprolactone and silane modification using organosilane (3-Aminopropyl)triethoxysilane (APTES). FTIR, XRD, HRTEM and water contact angle measurement manifests a successful organic covalent modification of HNT surface. The synergistic mechanical properties of covalently PCL modified HNT nanocomposite is because of ameliorated polymer-filler interface interaction that can be understood by the study of “Pukánszky model” and dynamic mechanical analysis (DMA). Rheology study of nanocomposites helped understanding the polymerfiller interaction. Good polymer-filler adhesion leads to a better state of filler dispersion in the polymer matrix as evident from the“effective free space length (Lf*)” of TEM analysis, WAXD, FESEM, and AFM study of nanocomposites.

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KEYWORDS: Nanocomposite, Interface adhesion, Microstructures, Electron microscopy 1.

INTRODUCTION Organic modification of various inorganic fillers is a commonly used technique to

increase its wettability of polymers. Halloysite nanotubes (HNT) are naturally occurring clay, primarily hollow tubular structure in the submicron range. Doping of polymers with HNT have recently got a tremendous research attention due to many salient features like improved mechanical, thermal, and fire-retardant performance

1-2

. An effective anchoring

of polymer chains within the lumens of HNT can improve polymer-filler interface adhesion of the HNT/polymer nanocomposites 3. HNT are silica based nanotubes that possess relatively lower hydroxyl density on their outer surfaces compared with other silica based inorganic fillers like fumed silica and montmorillonite (MMT). Therefore, the extent aggregation induced by the inter-tubular hydrogen bond type interaction between its surface hydroxyl groups is less. Hence, during the melt mixing the inter-tubular force of attraction is highly susceptible to the shearing action in the viscous polymer matrix. From the structural point of view, though HNT has similarity with CNTs, but unlike CNTs the HNT are non-toxic to living beings and the environment, and even it is biocompatible. Lately, increasing number of research groups are reported to have successfully fabricated and characterized HNT based polymer nanocomposites mainly focusing on mechanical properties 3-5. However, it is necessary to account for the available reports on the limitation of hydrophilic HNT dispersion in organic polymer matrix. Organophilization of inorganic fillers has got a lot of attention as a tool for improving polymer-filler affinity from researchers across the globe. The effect of surface modification on the mechanical


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properties of PP/HNT were investigated by Ning et al. 6. Guo et al. and HedickeHöchstötter et al. observed improvement in mechanical properties of PA6/m-HNT nanocomposites due to chemical surface modification of HNT 7-8. Chlorinated polyethylene (CPE) and ethylene methacrylate copolymer (EMA) are two widely used polymers in the field of cable, hose, pipe, and under the hood applications. The saturated ethylene backbone structure of both polymer results in high thermal stability, flame resistance property, chemical resistance, and excellent weatherability. Blending the two polymers in different composition caused appreciable improvement in mechanical properties and thermal stability as reported in an earlier publication by the same research group 9. CPE and EMA both have polar functional groups attached to the backbone, and hence a polymer-filler interaction can exist by the involvement of surface hydroxyl groups of HNT. But the polymer-filler interaction by the participation of C—Cl group of CPE and acrylate group of EMA with surface hydroxyl group of HNT may not be sufficient enough to form highly reinforced nanocomposite. This work reports a facile one step covalent modification of HNT by in-situ ring opening polymerization of ε-caprolactone and chemical grafting of APTES. The modified HNT was successfully dispersed in CPE/EMA (60/40 ratio) blend matrix by conventional melt mixing technique. Micromechanical modeling by “Pukánszky model” using the mechanical properties of prepared nanocomposites was carried out to find the degree of polymer-filler interface adhesion. The morphology study by TEM analysis was also used to determine the “Effective free space length (Lf*)” that represents the effect of modified HNT on its dispersion in the polymer matrix. A comparative study between PCL and



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APTES modified HNT has also been carried out to find the best method of HNT surface modification suitable for the CPE/EMA blend matrix. 2.

EXPERIMENTAL

2.1

Materials. Commercial grade CPE elastomer (CPE 360) with 36% Cl content,

having density of 1.213 g cm-3 was obtained from East Corp International, India. Commercial grade of EMA, Elvaloy®1330 with 30% methyl acrylate (MA) content was supplied by NICCO Corporation, Shyamnagar, India. Poly(ε-caprolactone) (PCL) with molecular weight (Mn) of 14,000, (3 Aminopropyl)triethoxysilane (APTES) and HNT with an average tube diameter of 55 nm and inner lumen diameter of 15 nm were procured from Sigma Aldrich. Catalyst Stannous octoate (Sn(Oct)2) and the monomer ε-caprolactone were purchased from Sigma Aldrich. Magnesium oxide (MgO) of density 3.58 g cm-3 was used as an acid scavenger for hydrochloric acid (HCl) produced during processing and molding. Dibutyltin dilaurate (DBTDL) and Irganox 1010® which were procured from Sigma Aldrich were used as the heat stabilizer of CPE and as antioxidant, respectively. 2.2.

Methods

2.2.1. HNT modification Covalent modification of HNT by ε-caprolactone. A chemical masterbatch of HNT was prepared by in-situ ring opening polymerization of ε-caprolactone monomers in HNT by bulk polymerization technique. First the HNT were swollen by ε-caprolactone monomer followed by ultrasonication for 30 mins in an inert atmosphere (N2). A dry toluene solution of catalyst stannous octoate Sn(Oct)2 was injected into the reaction chamber. A preevacuation treatment was carried out in order to remove the air present in the HNT lumen 4 ACS Paragon Plus Environment

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and thereby to facilitate the monomer and catalyst molecules enter into the nanotube lumen. Then the reaction was initiated by exposing the reaction chamber to 120 °C under nitrogen atmosphere. The reaction was stopped after 24 hrs by temperature quenching and then the resulting product was precipitated from cold methanol. It is named as PCL-g-HNT(C) and further it was recovered by filtration using THF. The obtained product was re-dispersed in THF and filtered at least for four times to ensure removal of any un-grafted polymer. Non-covalent modification of HNT by Poly(ε-caprolactone). The poly(ε-caprolactone) (PCL) flakes were dissolved in THF at room temperature by stirring for 12 hrs. HNT were dried in the vacuum oven at 70 °C for 8 hrs and then were dispersed in THF solvent by weighing required amount. Then the two mixtures are mixed together and stirred for another 24 hrs followed by ultrasonication. The prepared product named as PCL-g-HNT(P) was filtered with THF. The obtained product was re-dispersed in THF and was filtered, and the dispersion-filtration cycle was repeated at least for four times to ensure removal of ungrafted polymer. Covalent modification of HNT by APTES (Silane approach). Organosilane modified HNT was prepared by covalent grafting of APTES onto HNT surface by following the procedure usually applied for chemical grafting of silica-based fillers

10

. In a typical run, 4 mL of

APTES was dissolved in dry toluene. Approximately 10 g of HNT powder was added to the suspension that was dispersed ultrasonically for 30 mins. Evacuation pretreatment was carried out at this stage in order to remove air in the lumen and thereby to facilitate the APTES molecules enter the lumen of HNT. The suspension was then refluxed at 120 °C for 24 hrs under constant stirring in N2 environment. In the refluxing system, a calcium 5 ACS Paragon Plus Environment


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chloride drying tube was also attached to the open end of the condenser in order to ensure a dry environment. The resultant mixture named as APTES-g-HNT was filtered and extensively washed five times with fresh toluene to remove the un-grafted organosilane, and then dried overnight at 80 °C 11. 2.2.2. Preparation of C60E40 nanocomposites. CPE/EMA blend of 60/40 ratio was taken as the base polymer matrix for the nanocomposites. First, CPE was melted at 140 °C for two minutes along with the additives MgO, DBTDL, and Irganox 1010 which was followed by addition of EMA and mixing was continued for another 4 minutes. Then the synthesized PCL-g-HNT(C), PCL-g-HNT(P), APTES-g-HNT and pure HNT of equal quantity was added into the internal mixer and mixing was continued for another 6 minutes. The sample designations and their compositions are mentioned in Table 1. Sheets of 2 mm thickness of all mixes were prepared by compression molding. 2.2.3. Characterization techniques. Fourier transform infrared (FTIR) spectroscopy studies were performed on a Bruker Equinox 55 FTIR spectrophotometer, at a resolution of 2 cm−1, in the range of 4000–400 cm−1, and 64 scans were averaged out for each spectrum. X-ray diffraction (XRD) measurements were performed on X’Pert PRO, PANalytical instrument, with crystal monochromated CuKα radiation (λ=1.54 Å) in the angular range of 1–40° (2θ). The water contact angle was measured by a contact angle goniometer (RameHart instrument Co., Model 190F2). Following the initially proposed procedure by Rogers et al., a pressure of 30 MPa was applied to the powder samples with a metal plunger for 5 s at room temperature in order to obtain 3 mm thick compact clay disks 12. For analyzing the physicomechanical properties, the prepared nanocomposites were tested in a Hioks– 6 ACS Paragon Plus Environment

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Hounsfield Universal Testing Machine (Test Equipment, Ltd., Surrey, England) at a crosshead speed of 500 mm min-1, at room temperature. Dynamic mechanical analysis (DMA) of the neat CPE/EMA blend and its nanocomposites were determined in Dynamic Mechanical Analyzer DMA Q800 (TA Instruments, Lukens Drive, Newcastle, Delaware). The temperature sweep measurements were done under the tension mode in the temperature range from −75 °C to +100 °C at a heating rate of 3 °C min−1 with 0.1% strain and 1 Hz frequency. The strain sweep measurements in DMA were carried out at 25 °C over the range from 0.1% to 20% at a constant frequency of 1 Hz. The melt flow properties at steady and high shear rates of the neat C60E40 blend and its nanocomposites were evaluated by using fully computer-controlled Smart RHEO (CEAST S.p.A., Italy). The extrusions were performed at the processing temperature of 140 °C at eight different shear rates in the range of 10 to 2000 s− 1. As per the instrument specification the Bagley correction can be assumed to be negligible. The bulk morphology was analyzed by high-resolution transmission electron microscopy (HRTEM) using a JEM 2100 JEOL transmission electron microscope with a lanthanum hexaboride target, operating at 200 kV. The surface morphology of prepared nanocomposites were analyzed by using field emission scanning electron microscopy (FE-SEM) ZEISS MERLIN GEMINI 2®, operating at 20kV. The bulk morphology of the prepared HNT nanocomposite was also checked by slicing the compression molded samples using ultramicrotome machine. Intermittent contact mode atomic force microscopy, AFM (Agilent 5500 Scanning Probe Microscope) was used to carry out the investigation on morphology of the HNT nanocomposites. Thermal degradation study of neat blend and its nanocomposites were carried out using TGA



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STARe System, METTLER TOLEDO at a heating rate of 20 °C min-1 under nitrogen atmosphere over the temperature range of 35 °C to 600 °C. 3.

RESULTS AND DISCUSSION

3.1.

Characterization of pristine and modified HNT

3.1.1. WAXD and FTIR analysis. The covalent surface modification of HNT with PCL was carried out by following in-situ ring opening polymerization of ε-caprolactone using stannous octoate as catalyst. As mentioned earlier, HNT contains a sufficient amount of hydroxyl groups (OH) on its outer surface and inside the lumen wall. Such OH groups participate as active sites for ring opening of ε-caprolactone 13. The reaction was activated by adding catalyst stannous octoate to the reaction mixer at 120 °C. Here, the outer surface silanol hydroxyl groups (Si—OH) of HNT hollow tubes and the aluminols (AlOH) inside the lumen acts as initiator for polymerization which can be realized from the FTIR analysis of HNT, PCL-g-HNT(C) and PCL-g-HNT(P). The catalyst stannous octoate reacts with the OH groups and forms active sites for ring opening polymerization. For better understanding, a schematic of above mentioned process is shown in Figure 1-a. Covalent modification of HNT by organo-silane APTES was carried out according to the normally followed procedure for chemically grafting silica based fillers 11. As previously reported by Pasbakhsh et al., the OH group on outer surface and edge of the lumen wall of HNT reacts chemically with Si—OR groups of APTES which is presented schematically in Figure 1-b 14

. Figure 2-A shows the WAXD patterns of pristine HNT, PCL-g-HNT(P), PCL-g-

HNT(C), and APTES-g-HNT. For the unmodified and modified HNT samples, the observed peaks are close to the characteristic data of halloysite (JCPDS card No. 29-1487).


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Additionally, a small peak of 1 nm is observed in lower 2θ region. This peak corresponds to the small amount of hydrated particles in the HNT and is reported to be observed by S. Rooj et al. 15-16. Disappearance of this peak (1 nm) in modified HNT is most likely attributed to loss of interlayer water molecules. Two distinct diffraction peaks are observed for both the PCL-g-HNT at 2θ = 21.42 and 2θ = 23.73. These diffraction peaks were indexed to be (110) and (200) planes respectively; of an orthorhombic crystal structure of PCL and is an implication of PCL grafting on HNT surface

17

. Figure 2-B shows the baseline corrected

FTIR graphs of pure PCL, PCL-g-HNT(P), PCL-g-HNT(C), APTES-g-HNT, and pristine HNT. Here, the inner Al—OH groups that are in between the interface of the HNT walls are not reachable by ε-caprolactone monomer (as also observed from the Table S1). Therefore, its stretching vibration signals (3618 cm-1) are used as reference peak to evaluate the extent of modification. The reduced peak intensity of O—H- stretching vibration of inner surface hydroxyl groups at 3696 cm-1

in modified HNT and slight shifting of

perpendicular Si—O stretching peak from 689 and 753 cm-1 in HNT towards lower wavenumber (along with reduced intensity) indicates effective grafting. This procedure shows that the extent of modification follows an order of: PCL-g-HNT(C)> APTES-gHNT> PCL-g-HNT(P)

11

. Furthermore, appearance of a sharp carbonyl peak at around

1725 cm-1 in case of PCL modified HNT as shown in the highlighted region in Figure 2-B, is a clear intimation of successful covalent grafting of PCL onto the HNT surface. Figure 2B also demonstrates the evolution of two tiny additional peaks of chemical group C—H2deformation at 1494 cm-1 and N—H2- deformation (scissoring) at 1556 cm-1 in organosilane modified HNT. These two peaks are not present in unmodified HNT

18

. The

frequency and assignment of each vibrational mode observed are enlisted in Table 2 19. A 9 ACS Paragon Plus Environment


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very intense peak at 1048 cm-1 was owing to stretching vibration of Si—O—Si from silicate 20. However, the characteristic band at 1048 cm-1 for Si—O—Si stretching in PCLg-HNT was not so intense. This is because the polymer has multiple bands in the same spectral range and by far PCL is attached in large content to HNT surface 21. 3.1.2. TEM observation. The fine nanostructures of pristine HNT, PCL-g-HNT(P), PCLg-HNT(C), and APTES-g-HNT samples were measured by high-resolution transmission electron microscopy (HRTEM). The representative HRTEM images of unmodified and modified HNT are shown in Figure 3. Figure 3-a, b, and c show the HRTEM images of pristine HNT in three different locations including magnified images in Figure 3-b and c. The halloysite nanoparticles have a cylindrical shape containing clearly visible walls and a transparent central area that runs longitudinally along the open-ended hollow cylinder in Figure 3-a. Image J® (NIH, USA) software was used to calculate the dimensions of HNT. The outer diameters of pristine HNT are found to be in the range of 55 to 62 nm, and the inner lumen diameter is about 18 to 22 nm. The spiral structural layers of HNT walls are clearly visible in the Figure 3-b and c where the nanotube diameters are congruent with the results obtained in Figure 3-a. Averages of 50 individual nanotubes at high magnification from the different area of each sample are taken into consideration for the measurement of the diameter, so as to estimate the modification of HNT qualitatively. Figure 3-d, e, and f depict the representative HRTEM images of PCL-g-HNT(P), PCL-g-HNT(C), and APTESg-HNT respectively in order to have an idea about an essence of the modification. The individual nanotubes of physically grafted HNT i.e. PCL-g-HNT(P) (Figure 3-d) are morphologically same as pristine HNT, and the diameter showed no difference in values


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from pristine HNT. In the case of PCL-g-HNT(C) in Figure 3-e, an uneven layer of polymer was observed on the outer surface of HNT. It is also clearly evident that the measured diameter of the HNT is 65-70 nm, which is appreciably larger than those of pristine HNT. The increase in diameter after covalent grafting can be attributed to a large amount of PCL chains wrapped all over the nanotube surfaces as manifested by the FTIR analysis. Hence, the TEM morphology suggests that in-situ ring opening polymerization of ε-caprolactone is pretty effective for functionalization of HNT

22

. Similarly, the organo-

silane modified HNT i.e. APTES-g-HNT also showed a marginal increase in the outer diameter of HNT with prominent surface roughness as presented in Figure 3-f. As proposed by Yuan et al., there is a chance for the organo-silanes (APTES) to oligomerize or polymerize. Further, the oligomerized APTES can covalently react with the grafted APTES to form a cross-linked network 11, 16. 3.1.3. Contact angle measurement. Contact angle measurement is one of the conventional techniques to measure the hydrophobicity of organo-modified inorganic fillers. The higher contact angle value means better hydrophobicity, which in turn directly reflects the high degree of grafting. A goniometer was used to measure the contact angles of the water on pristine HNT and modified HNT as shown in Figure 4. It is clear that when de-ionized water droplet (1.0 µL) was dropped onto the surface of pristine HNT, a very flat droplet formed on its surface with a static angle of c.a. 27.5° (Figure 4-a). On the contrary, all the modified HNT show higher values of static contact angle. The degree of hydrophobicity increased upon organic modification of HNT and it follows the increasing order of: HNT (27.5°) < PCL-g-HNT(P) (42.0°) < APTES-g-HNT (52.2°) < PCL-g-HNT (C) (56.0°).



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From the contact angle data, the interfacial adhesion energy (or work of adhesion) between the solid and the liquid phases can also be calculated by Young's equation 23:

Wa = γ L (1+ cosθ )

(1)

Where, the γL is the surface energy of pure water (72.8 mJ m-2

24

), and θ represents the

contact angle values. The calculated interfacial adhesion (Wa) shown in Figure 4 reveals highest value in case of covalently modified PCL-g-HNT(C)

25

. The result presented in

Figure 4 fully corroborates with HRTEM results as discussed in the preceding section. Hence, the contact angle measurement claims that in-situ ring opening polymerization of εcaprolactone caused higher degree of PCL chain grafting onto the HNT surface as compared to APTES silane grafting. This efficient grafting of PCL onto HNT further led to an increase in the hydrophobicity and wettability in polymer matrix. 3.2.

Characterization of nanocomposites

3.2.1. Mechanical properties. The bar diagram of the ultimate tensile strength, tensile modulus at 100% elongation, and % elongation at break values of unmodified and modified HNT nanocomposites are shown in Figure 5. In this case, the neat C60E40 base blend matrix is also taken into consideration in order to depict a clearer picture of the effect of unmodified and modified HNT on CPE/EMA blend. Special emphasis has been put on the ultimate tensile strength property of the prepared nanocomposites by mentioning the % change in tensile strength over neat blend. It was observed that the ultimate tensile properties followed the expected trend of higher values in the covalently modified HNT based nanocomposites compared to (nano)composites of pristine HNT. From the tensile data, it is clear that the covalently modified HNT based nanocomposites displays higher


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ultimate tensile strength compared to non-covalently modified HNT and pristine HNT based (nano)composites. In addition, the nanocomposite C60E40/PCL-g-HNT(C) showed superior tensile strength over C60E40/APTES-g-HNT. The inclusion of HNT in both pristine and modified form increased the tensile strength of the C60E40 blend. The addition of covalently modified PCL-g-HNT caused a remarkable 44.36% increment in ultimate tensile strength value while APTES-g-HNT showed 15.45% increase. On the other hand, non-covalently modified C60E40/PCL-g-HNT(P) and pristine C60E40/HNT showed only 11.86% and 10.58% hike in tensile strength values. The reason for such appreciable reinforcement of HNT in C60E40/PCL-g-HNT(C) can apparently be attributed to the strong polymer-filler interaction with better state of dispersion in polymer matrix. The successful grafting of organic PCL polymer chains onto the HNT surface not only enhanced the hydrophobicity of HNT but also amended the polymer-filler interface interaction by acting as a compatibilizer between HNT and CPE/EMA blend. In addition, the high mechanical shearing in viscous polymer melt during melt mixing process has a profound effect on the quality of PCL modified HNT dispersion. An exquisite adhesion of covalently modified HNT with the base polymer matrix helps transfer the applied tensile stress along the polymer-filler interface expeditiously. A very good dispersion of modified HNT in the polymer matrix naturally has enormous tendency to show synergistic mechanical properties. Figure 5 also demonstrates an increasing trend in elongation at break (%) on addition of HNT that is because of entrapped polymer chains inside the HNT lumen as previously reported by Ismail et al. and Raman et al.

16, 26

. The in-situ

polymerization of ε-caprolactone using hydroxyl groups of HNT as reactive site for ring opening causes grafting of PCL chains even inside the lumen surface. Therefore, the 13 ACS Paragon Plus Environment


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amount of polymer chains entering the HNT lumen during melt mixing is expected to be much more in C60E40/PCL-g-HNT(C). During tensile testing, these trapped polymer chains slowly arise from the lumen and thereby leading the nanocomposites to higher resistance to break16. 3.2.2. Micromechanical modeling. The ultimate tensile strength of a polymer composite is widely determined by strain hardening due to polymer chain orientation and polymerfiller interaction. According to Interfacial adhesion reinforcing theory as mentioned by Liang 27, the higher the interfacial adhesion strength, the better is the reinforcement of the polymer composites. Moreover, under the same interfacial adhesion strength, the larger the interfacial area of the fillers, the better is the reinforcement of the polymer composites. A semi-empirical equation on the effect of spherical fillers in ultimate tensile properties of polymer composites was proposed by Pukánszky

28

. But recently, the model has

successfully been applied to anisotropic fillers such as sepiolite the needle-like nanoparticles, multiwalled carbon nanotubes (MWCNTs), wood fibers, and nano hydroxyapatite (nHA) nanorods

29-32

. The increase in material strength due to polymer

chain orientation along the direction of tension is called “strain hardening” which can be expressed as:

σ T = λn Where

σT =λ•σ

and

(2)

λ =

L L0

(3)

In this case, σT is the true stress, and σ is engineering stress; whereas, λ is the relative elongations while the L0 and L are the original and actual lengths of specimen. The “n” is a


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parameter characterizing the strain hardening tendency that can be calculated by a curvefitting procedure. In polymer composites, when polymer adheres to the surface of fillers, the tensile strength increases. The expression adapted for tensile yield stress can also be adjusted for tensile strength:

σ Tred = σ T 0 • exp (B φ f

)

(4)

Where, σTred is reduced true tensile strength, σT0 is the true tensile strength of polymer matrix and B is polymer-filler interfacial interaction parameter. Combining, Equation (2) – (4) gives the final expression: σ

T

B =

= σ

1

φ

f

T 0

 (1 − φ f )  • λn   exp  (1 + 2 . 5 φ f )

 1 • log  n  λ

 σ T  σ T0

  1 + 2 .5φ     1 − φ f

f

(B φ )

(5)

   

(6)

f

Pukánszky worked with various polymers, and the results of his work reflect that the strain hardening behavior of the polymer matrix is independent of filling 28. Here, the value of “n” can be determined from the plot of log σT vs. log λ as shown in Figure 6. The “n” value in case of CPE/EMA polymer matrix is found to be 1.65 from the slope of the plot in Figure 6. B is an empirical parameter that can be determined for all the prepared nanocomposites by putting the experimental results in Equation (6). In this case, it is assumed that interfacial interaction is the only responsible factor for the observed differences in properties of the nanocomposites. As a basic line of thought, when polymer adheres to the filler, an interface gets created in between the polymer and filler. The property of this interface naturally dictates the overall mechanical properties of the composite. The calculated B value obtained from the Pukánszky’s model of Equation (6) for all the prepared nanocomposites



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are tabulated in Table 3. If B value is less than 1, the interfacial adhesion will be weak; whereas, if B value is closer to 3, then the interfacial adhesion between polymer-filler will be strong enough to transfer the applied stress effectively 27, 32. It can be observed that the B values of all prepared nanocomposites with equal filler fraction are higher than 2. The covalent grafting of HNT led to increment in the B values of corresponding nanocomposites. Moreover, PCL-g-HNT(C) showed highest B value of 3.36, suggesting good interfacial adhesion between the polymer matrix and PCL modified HNT. The effect of unmodified and modified HNT on the viscoelastic and other rheological properties of C60E40 blend matrix are studied in the subsequent sections. The morphology of polymer nanocomposites has an enormous effect on its mechanical properties that has to be taken into account for the study of its structure-property relationship. A comprehensive morphology study of the prepared nanocomposites has been made in order to understand its effect on observed mechanical properties.

3.2.3. Dynamic mechanical analysis (DMA). DMA is one of the most widely accepted techniques to investigate the viscoelastic behaviour of polymeric materials and to evaluate the effect of filler addition on polymer nanocomposites

33-34

. In this regard, dynamic

mechanical properties of polymeric material exhibit much practical significance when tested over a temperature range 35. Figure 7-a and b presents tanδ and the storage modulus of neat C60E40 blend and its prepared HNT nanocomposites over a temperature range of 75 °C to 100 °C. The temperature corresponding to the maximum of tanδ in the temperature sweep plot of neat blend represents the glass-rubber transition temperature (Tg) of the polymer matrix. In case of filled polymer system, dynamic mechanical testing is frequently


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used to study the effect of various fillers on the Tg of polymers 36. Table 4 summarizes the Tg of all the prepared samples along with its storage modulus above Tg (100 °C). From the Table 4, it is clear that all nanocomposites showed higher Tg value compared to the neat C60E40 blend matrix. The pristine HNT and non-covalently modified HNT based nanocomposites (i.e. C60E40/HNT and C60E40/PCL-g-HNT(P) respectively) did not show appreciable shifting of Tg. This insignificant change in Tg of the C60E40/HNT and C60E40/PCL-g-HNT(P) nanocomposites implies that these fillers have little effect in reinforcing the base polymer matrix

37

. The reason can be explained on the basis of poor

polymer-filler interaction between the filler and the C60E40 blend matrix. The hydrophilic nature of the pristine HNT and PCL-g-HNT(P) surface barely provides any scope for strong chemical or physical interaction with C60E40 blend matrix. The damping (height of the tanδ peak) indicates the ability of a material for dissipating the applied energy. At Tg region, the long range macromolecular chains attain mobility by dissipating energy through viscous movement 38. The C60E40/HNT and C60E40/PCL-g-HNT(P) nanocomposites also indicated a slight reduction in tan δ peak height, which refers to little reinforcing effect of these fillers (as also can be realized from the ultimate tensile data). On the contrary, the substantial increase in Tg of the covalently modified HNT based nanocomposites (i.e. C60E40/APTES-g-HNT and C60E40/PCL-g-HNT(C)) along with distinct peak broadening suggests good polymer-filler interfacial adhesion

39

. Furthermore, the C60E40/PCL-g-

HNT(C) nanocomposite displayed highest Tg value with minimum tanδ peak height, which is due to the decreased mobility of the polymer chains in the vicinity of the fillers. The existence of good polymer-filler interfacial adhesion in C60E40/PCL-g-HNT(C) nanocomposite is expected to be due to higher degree of modification of HNT by PCL



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polymers (as discussed in the HRTEM and water contact angle study). Homogeneous filler dispersion with good polymer-filler interfacial interaction in nanocomposite is more likely to increase the amount of polymer chains with a restricted movement that in turn increases the Tg and reduces the tanδ peak height

40

. Figure 7-b shows the variation of storage

modulus as a function of temperature for C60E40 blend and its nanocomposites. From the Figure 7-b, it can be understood that below Tg, the E’ value of both unfilled and filled samples are very high due to the molecular stiffness. But no significant variation was observed in the E’ values of neat C60E40 blend and its nanocomposites in the glassy region. It is worth to mention that storage modulus E’ of polymers in the glassy region is primarily dictated by the strength of the intermolecular force and the way of packing of polymer chains. At high temperature i.e. above Tg, the E’ value is determined by the much compliant amorphous regions

37

. The values of E’ at high temperature i.e. above Tg (100

°C) are mentioned in Table 4. All the nanocomposites showed higher E’ value over the neat C60E40 blend at the rubbery plateau as can be realized from the data of Table 4. This observation can be ascribed to the reinforcement mechanism of fillers, as these fillers have a larger effect in raising the modulus above Tg

34, 41

. Besides, the nanocomposites

underwent lesser drop in the storage modulus in comparison to the neat blend on passing across the glass-rubber transition temperature and the trend is in the follows order: C60E40/HNT < C60E40/PCL-g-HNT(P) < C60E40/APTES-g-HNT < C60E40/PCL-gHNT(C). Effective covalent surface modification of HNT by organic APTES and PCL increased the filler wettability in the polymer matrix by improving polymer-filler interfacial adhesion. During melt mixing of viscous polymer, good polymer-filler interfacial interaction leads to homogeneity in the dispersion of filler under severe mechanical shear.


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The efficient covalent grafting of HNT by PCL over APTES unanimously maximized interfacial adhesion of PCL-g-HNT(C) with the polymer matrix and thereby caused the formation of immobilized regions around the filler surface 42. Moreover, the PCL polymer chains provided an added advantage in improving compatibility with CPE/EMA blend matrix. Hence, covalent modification of HNT by PCL is highly anticipated to be an effective technique for improving polymer-filler interfacial adhesion in C60E40 matrix. The above study of the viscoelastic property of polymer nanocomposite depicts a clearer picture of reinforcement by the covalently modified HNT.

3.2.4. Rheology study of nanocomposites DMA strain sweep. In order to investigate the polymer-filler and filler-filler interaction of unmodified and modified HNT based nanocomposites of the C60E40 blend, strain sweep measurements were performed

43-46

. The effect of pristine and modified HNT loading on

storage modulus E’ of nanocomposites at dynamic strain amplitude is documented in Figure 8. All samples showed highest storage modulus E’ value at low strain amplitude, and the value gradually decrease as the strain amplitude increases. Such observation is due to the reason that at a very low strain the material gets plenty of time to undergo molecular rearrangement so as to minimize the local stress than at high strain

47

. If the filler-filler

interaction is stronger than the polymer-filler interfacial interaction, then a large filler-filler network exists in the filled system. In that case, the drop in storage modulus E’ with increasing strain amplitude is much sharper than the unfilled system. This behaviour of filled polymer system is commonly known as “Payne effect” that was recently also been reported for low filler loading by Das et al. and Rooj et al.

48-50


. The magnitude of “Payne


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effect” can be represented as (E’0 — E’∝); where E’0 and E’∝ refer storage modulus value at low and high strain amplitude. The greater significance of “Payne effect” in a highly filled system is due to the disruption of filler-filler network structures with increasing strain amplitude. All nanocomposites showed a slightly higher value of storage modulus E’ at low strain compared to the neat C60E40 blend matrix. This observation is naturally due to the presence of rigid HNT fillers in low concentration (5 wt%). A significant increase in storage modulus E’ value of nanocomposites at high strain % over the neat blend complies with the reinforcing nature of the fillers from the ultimate tensile data

44

. The remarkably

low “Payne effect” in covalently modified HNT based nanocomposites (as presented in the inset table of Figure 8) implies an existence of good polymer-filler interfacial interaction with lesser filler-filler aggregate in the polymer matrix. Taking the absolute E’ value and “Payne effect” as the measure of reinforcement and filler-filler network structure respectively, the used fillers are ranked according to their reinforcing efficiency as: pristine HNT < PCL-g-HNT(P) ≤ APTES-g-HNT < PCL-g-HNT(C). This progression derived from the rheological data of strain amplitude sweep of nanocomposites well corroborates with the former mechanical property observation.

Melt Rheology study. The apparent shear viscosity vs. shear rate plots for the neat C60E40 blend and its unmodified and modified HNT based nanocomposites were plotted in Figure 9. The decreased apparent shear viscosity with increasing shear rate indicates pseudoplastic or shear thinning behaviour of all studied samples. The prepared nanocomposites showed higher viscosity than the neat blend system

51

. Since, all the prepared nanocomposites

contained equal filler concentration (5 wt%); therefore a comparative study on the effect of


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HNT modification in melt shear viscosity can be made. The small increase in viscosity of non-covalently modified HNT and pristine HNT based nanocomposites over neat C60E40 blend is due to the presence of occluded rubber in the locality of aggregated HNT fillers 52. However, compared to the neat C60E40 blend, both the covalently modified HNT based nanocomposites i.e. C60E40/APTES-g-HNT and C60E40/PCL-g-HNT(C) showed much increase in viscosity, which is quite expected from the previous studies. This is an indication of good polymer-filler interfacial adhesion and uniform dispersion of fillers in the polymer matrix that induces larger resistance to flow and deformation compared to the C60E40/HNT and C60E40/PCL-g-HNT(P) nanocomposites

53

. A similar explanation of

higher shear viscosity of PEO/Clay intercalated nanocomposite over its immiscible nanocomposite was reported by Choi et al. 54.

3.2.5. Morphology study of nanocomposites Wide angle X-ray diffraction (WAXD) analysis. The WAXD patterns of the prepared nanocomposites of pristine HNT and modified HNT are shown in Figure 10 including the diffractogram of pure HNT in the inset. The diffractogram of the C60E40/HNT nanocomposite displays an intense peak at 11.96° with basal spacing of 0.739 nm for the (001) plane of HNT. This peak of HNT has shifted to lower 2θ value in case of C60E40/PCL-g-HNT(C) (2θ = 11.60°, d-spacing of 0.761 nm) and C60E40/APTES-g-HNT (2θ = 11.68°, d-spacing of 0.750 nm); but no such shifting is observed in case of C60E40/PCL-g-HNT(P). This reduction in 2θ value and increase in d-spacing of the HNT in the covalently modified nanocomposites may be attributed to the intercalation of HNT by other materials like polymer chains and MgO. This supposition is in accord with the



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results of Yuan et al. and Rooj et al.

11, 55

. As represented in the Figure 10, disappearance

the peak at 20.06° of pure HNT (as shown in inset) in WAXD patterns of nanocomposites is considered as an evidence for good dispersion of the HNT 55. In addition to that, shifting of the 2θ value to lower side of some other characteristics peaks of nanocomposites in Table S2 substantiate successful covalent modification of HNT.

Transmission electron microscopy (TEM) analysis. As a rule of thumb, the state of nanoparticle dispersion in the polymer matrix is a powerful, decisive factor in determining ultimate properties of nanocomposites. Uniform and homogeneous dispersion of nanoparticles can lead the mechanical property of nanocomposites to very high level provided with the condition of high degree of polymer-filler interfacial adhesion. An effective organic surface modification of inorganic filler increases the wettability of nanoparticles and hence its dispersion in the polymer matrix. On the contrary, the aggregated microstructures of unmodified hydrophilic nanoparticles in the polymer matrix act as the point of stress–concentration, which leads to deterioration of mechanical properties. Figure 11 demonstrated the state of dispersion of the pristine and modified HNT into the CPE/EMA polymer matrix at low magnification of 0.2 µm. The bright background shows the polymer matrix while the dark tube-like structures are individual HNT. It can be observed that individual HNT nanofillers are accompanied by some larger aggregates in certain nanocomposites. Aggregates are only qualitatively defined as ensembles of the nanoparticles. In the case of C60E40/HNT (nano)composites in Figure 11-a, large number of aggregates and agglomerates of HNT are observed. Presence of large aggregates of HNT is an indication of poor quality dispersion in C60E40/HNT (nano)composite. Almost


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similar type of morphology was observed in the case of non-covalently grafted HNT based nanocomposite (C60E40/PCL-g-HNT(P) in Figure 11-b) with lesser number of aggregates. On the contrary, in the case of nanocomposites of covalently modified HNT i.e. PCL-gHNT(C) and APTES-g-HNT showed better state of dispersion with fewer agglomerates. Here individual nanotubes are apparently well separated and dispersed from each other. Indeed, individual nanotubes without any remaining bundle or aggregate are observed in the case of in-situ covalently modified HNT based nanocomposites (i.e. C60E40/PCL-gHNT(C)) in Figure 11-d. The treatment for calculating the agglomeration length (La) and free-space length (Lf) was done by using image analysis software (Image J 1.6r). A quantitative approach was employed to show the state of nanofiller dispersion in TPU matrix. In nanocomposites, as the aggregates become larger relative to the free-space, they turn out to be the most favorable location for failure which reflects in macroscopically diminished properties. A higher value of effective free-space length Lf* represents better state of dispersion at the similar filler loading in the nanocomposites. The effective freespace length Lf*, is the equivalent free-space length for infinitely small agglomerations and Lf* is defined as 56, n     L 1  a  L f * = L f 1 +  ⋅ a Lf      

(7)

Equation (7) can be rearranged by

1  L f * = L f +  La n × L f 1− n  a 


(8)


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Where La is the agglomeration length, Lf is the free-space length, “a” is the critical ratio of agglomeration length to free-space length, and “n” is a sensitivity exponent. Generally, the modified free-space length is insensitive to agglomeration size as long as the agglomerations are smaller than the free-space length. Using the sensitivity exponent n = 2 in order to avoid some discrepancies as reported by Khare et al., the efficient free space length was measured for all prepared nanocomposites reference

56-57

. The distance with

zero gray value corresponds to the matrix region separating two individual HNT separated at a minimum distance, and it is a measure of the free-space length (Lf). The calculated effective free-space length Lf* value of all prepared nanocomposites are tabulated in Table 3. This observation can be well explained on the basis of the results obtained from the morphology study of unmodified and organo-modified HNT. The monolithic surface coating of HNT by PCL chains in PCL-g-HNT(C) and by organo-silane (APTES) in APTES-g-HNT resulted in high degree of hydrophobicity (as obvious from the contact angle measurement). The amended hydrophobicity of covalently modified HNT further increased its wettability by polymer matrix. Consequently, the state of dispersion of HNT in C60E40/PCL-g-HNT(C) and C60E40/APTES-g-HNT is much better than C60E40/HNT (nano)composites. For further comparison, TEM image for non-covalently modified HNT based nanocomposites i.e. C60E40/PCL-g-HNT(P) with equal filler fraction is shown in Figure 11-b. Here, bundle of nanotubes are much more prominent compared to individual tubes. Existence of larger quantity of HNT bundles in C60E40/PCL-g-HNT(P) clearly suggests the benefit of in-situ covalent modification of HNT over those of ex-situ (noncovalent) method. Hence, the covalent modification of HNT using PCL is the most efficient method for causing better deagglomeration by efficient hydrophobization HNT. Moreover,


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the nanocomposite C60E40/PCL-g-HNT(C) showed even better dispersion compared to C60E40/APTES-g-HNT that is because of more efficient grafting of HNT with PCL chains than APTES molecules. This observation is clearly assessable also from the contact angle measurement and HRTEM images of modified HNT. Hence, an efficient and successful grafting of PCL caused a desirably good dispersion of HNT in the polymer matrix. This explanation is manifested in the greater improvement of mechanical properties, which is discussed in the preceding section. Hence, the overall morphology study of C60E40/PCLg-HNT(C) is in well corroboration with the micromechanical modeling based on the experimental tensile data and the rheology study. The high polymer-filler interfacial adhesion in the nanocomposites along with good filler dispersion in the matrix caused a considerable improvement in mechanical property.

Field emission scanning electron microscopy (FE-SEM) study. For the better understanding of morphology, the FE-SEM and TEM images of unmodified and modified HNT nanocomposites at high magnification are shown in Figure 12. The Figure 12-a, b, c, and d depicts the FE-SEM images of pristine HNT based (nano)composite i.e. C60E40/HNT, C60E40/PCL-g-HNT(P),

C60E40/APTES-g-HNT, and C60E40/PCL-g-

HNT(C) respectively. Whereas, the Figure 12-e, f, g, and h represents the corresponding TEM images of C60E40/HNT, C60E40/PCL-g-HNT(P),

C60E40/APTES-g-HNT, and

C60E40/PCL-g-HNT(C) respectively. No momentous changes were observed in the appearance of HNT in C60E40/HNT (nano)composite except for the large number of aggregates (encircled with red color). Whereas, in case of modified HNT based nanocomposites, some visible changes were observed in the appearance of HNT in the



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polymer matrix. Considerable surface roughness was also observed for the modified HNT based nanocomposites as it can be clearly seen from both the FE-SEM and TEM images. The HNT walls were detected to be ruptured and thereby form some clay like features as clearly visible in Figure 12-b, c, d and f, g, h respectively. Such clay like gestalt and individually separated nanotubes in both FE-SEM and TEM photomicrographs are indicated with blue color arrow marks in Figure 12. The noticeable change in HNT morphology in the nanocomposites is fundamentally speculated to be because of better polymer-filler interaction. It is noteworthy to mention that the addition of reinforcing filler in the polymer matrix causes increase in overall viscosity that is solely because of good polymer-filler adhesion 58-59. In general, polymer chains firmly connect to reinforcing fillers either by chemical bonding or secondary force of interactions (like hydrogen bonding) or via. physical adsorption as in the case of carbon black. In this case, the successful hydrophobization of HNT by covalent modification, the filler–filler interaction is lowered. As referred by Fröhlich et al., the thin layers of polymer between the filler particles have restricted chain movement due to polymer-filler interaction

60

. Lowering of filler-filler

interaction and increased polymer-filler interaction caused enhancement in viscosity of the system

61

. This rise in viscosity in turn induces the filler to undergo massive shearing and

heat build up during mixing. The reinforcing fillers experience severe mechanical shear in the polymer melt (like elastomers) during processing. The non-covalently modified HNT based nanocomposite C60E40/PCL-g-HNT(P) in Figure 12-b and f showed little modification in its morphology with fewer agglomerates than C60E40/HNT. On the other hand, the covalently modified HNT showed appreciable difference in its appearance in C60E40/APTES-g-HNT and C60E40/PCL-g-HNT(C) nanocomposites (Figure 12-c, d and


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g, h respectively). Also, here the individual HNT are well separated from each other with lesser number of small aggregates. Though the exact mechanism of such morphology change in HNT is not clear, further detail study is necessary to have an insight into it. Nevertheless, such morphology change in modified HNT nanocomposites has proved to be expedient in terms of improvising mechanical properties of nanocomposites.

Atomic force microscopy (AFM) study. In order to fortify the results obtained from HRTEM and FE-SEM analysis of nanocomposites, AFM study was also carried out which represents much larger region. Though AFM is a technique to check the surface topography or morphology, but here the bulk morphology of the prepared nanocomposites was checked by slicing the compression molded nanocomposite samples by ultramicrotomy. The 2D phase images of AFM of all the four nanocomposites are shown in Figure 13 in order to check the type of dispersion of nanoparticles in the polymer matrix. The image of (nano)composite (Figure 13-a) of pristine HNT (C60E40/HNT) clearly reveals a significant amount of HNT agglomerates. The average width of the agglomerates is found to be ranging from 300 nm to 900 nm, distributed all over the polymer matrix. While the Figure 13-b shows the presence of HNT bundles in C60E40/PCL-g-HNT(P) that are much smaller than the agglomerates found in pristine HNT (nano)composites. However, the presence of individual HNT all along the surface were found in C60E40/APTES-g-HNT and C60E40/PCL-g-HNT(C) nanocomposite as shown in Figure 13-c and d. Though some of them are in associated form but many proved to be individually separated and spread all over the surface. Notably in C60E40/PCL-g-HNT(C) nanocomposite, the individual HNT count is prominent over its agglomerates. Indeed, a negligibly low level of agglomerates of



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HNT is found as clearly understood from the Figure 13-d. This observation is in well agreement with the representative HRTEM results of nanocomposites. The magnified image in 1µm scale of C60E40/PCL-g-HNT(C) in Figure 13-d is shown with an arrow mark, where the profile of individual nanotube represented a length of ca. 107 nm of a single HNT. The good dispersion of covalently modified HNT in the polymer matrix is certainly due to higher polymer-filler interfacial adhesion. Such high polymer-filler interaction is a direct consequence of effective grafting of PCL and APTES onto HNT surface (as evident from the HRTEM and XRD study of HNT). However, the extent of PCL chain grafting onto HNT surface being better than APTES resulted better deagglomeration and dispersion of HNT in the polymer matrix.

3.2.6. Thermogravimetric analysis (TGA). The inclusion of inorganic nanofillers has both positive and negative effect on the thermal stability of polymer matrix. The thermal stability of polymer nanocomposites is mainly dependent on the chemical structure of polymer matrix, polymer-filler interaction, and the morphology of filler inside the matrix. In order to investigate the effect of pristine and modified HNT on the thermal stability of C60E40 blend, TGA study has been carried out in an inert atmosphere of N2. The TGA and the derivative curves of neat C60E40 blend and its (nano)composites are plotted in Figure 14-a and b respectively. The Figure 14 displays two-step degradation behavior of neat C60E40 blend and its nanocomposites. The first step is associated with the evolution of hydrochloric acid (HCl) due to decomposition of CPE. The second step corresponds to degradation of polyenes of CPE and EMA copolymer. From the TGA plot, the onset degradation temperature Ti, (corresponding to 5 wt% loss) and T50 (temperature


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corresponding to 50 wt% loss), the maximum weight loss temperatures of the mentioned two steps (T1max and T2max) along with their char yield content (wt%) are summarized in Table S3 of supporting information. As shown in Table S3, the Ti of the neat C60E40 blend is 293.2 °C and the value increased on addition of pristine and modified HNT in small concentration (5 wt%). The addition of pristine HNT and PCL-g-HNT(P) caused an improvement in Ti value of around 4.9 °C and 6.4 °C over the C60E40 blend respectively. However, this temperature for covalently modified HNT based nanocomposites i.e. C60E40/APTES-g-HNT and C60E40/PCL-g-HNT(C) are 14.3 °C and 24.9 °C higher than neat C60E40 blend respectively. Further, the T50%, and the residue content also followed the similar trend. Besides, the higher thermal stability of covalently modified HNT based nanocomposites is mainly the result of the formed char that hinders the diffusion of volatile products by creating a tortuous path known as “Labyrinth effect”

62-63

. According to

Gilman, the barrier property (both thermal barrier and mass transport barrier) slows down the escape of the volatiles at the initial degradation and this effect is prominent in case of plate like clay structures 64. Therefore, the barrier effect of HNT is lesser than nanoclay due to its cylindrical structure. However, Du et al. reported that the lumen of HNT is capable of entrapping the degradation products of C60E40 blend matrix. This phenomena eventually results in effective delay of the mass transport and thereby caused a remarkable increase in thermal stability 65. The data from the Table S3 clearly indicates a highest thermal stability of C60E40/PCL-g-HNT(C) nanocomposites over the neat blend. It is reported in the literature that nanocomposites with a better state of filler dispersion generally lead to a higher thermal stability over the conventional micro-composites 66-67. The even distribution of PCL modified HNT in the C60E40 matrix (as clearly observable from the morphology



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study of nanocomposites) resulted in higher randomness of lumen ends. This randomness leads to enhanced effectiveness in entrapment of degradation products. Therefore, nanocomposites of PCL modified HNT has significantly high thermal stability than the other nanocomposites

65

. Also, the thermal stability of nanocomposites not only is

dependent on filler dispersion but also related to polymer-filler interfacial interaction

68-69

.

The successful covalent modification of HNT by APTES and PCL (as evident from the FTIR, HRTEM, and Contact angle data) organophilized the hydrophilic HNT and thereby reduce the filler-filler interaction. In this case, the grafted PCL chain plays a significant role in improving the polymer-filler interaction with CPE/EMA blend matrix that can be realized from the highest B value of C60E40/PCL-g-HNT(C) obtained from the Pukánszky’s micromechanical model.

4.

CONCLUSION The covalent organic modification of HNT by in-situ ring opening polymerization

of ε-caprolactone and direct chemical grafting of (3-Aminopropyl)triethoxysilane (APTES) was confirmed with the help of FTIR, XRD, HRTEM, and water contact angle measurement. The ultimate tensile properties of nanocomposite C60E40/PCL-g-HNT(C) and C6E40/APTES-g-HNT showed an appreciable increment in comparison to pristine and non-covalently modified HNT based nanocomposites. The synergistic mechanical properties of C60E40/PCL-g-HNT(C) over C6E40/APTES-g-HNT is anticipated to be because of high polymer-filler interfacial adhesion due to competent grafting of PCL polymer chains onto HNT (as evident from the TEM and water contact angle data). The morphology study using TEM, FE-SEM, and AFM proposes a better state of filler


Page 31 of 51

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dispersion

in

C60E40/PCL-g-HNT(C),

and

C6E40/APTES-g-HNT

compared

to

C60E40/HNT and C60E40/PCL-g-HNT(P). Moreover, the micromechanical modeling of “Pukánszky model” based on experimental data well corroborates with morphology results revealing highest interface interaction parameter B value for C60E40/PCL-g-HNT(C). The temperature sweep study in DMA showed a discernible increase in Tg values of covalently modified HNT based nanocomposites over the pristine HNT and non-covalently modified HNT based nanocomposites. The maximal value of Tg with an apparent broadening in tanδ peak and superior storage modulus of C60E40/PCL-g-HNT(C) at 100 °C indicates highest polymer-filler interaction in the system. Hence, the results from ultimate tensile tests are well supported by the dynamic mechanical study of the nanocomposites. In strain sweep test, substantial increase in storage modulus E’ value of all the nanocomposites of HNT at high strain % over the neat blend appeases the claim of reinforcing nature of the fillers from the ultimate tensile data. The remarkably low “Payne effect” in C60E40/PCL-g-HNT(C) implies an existence of good polymer-filler interfacial interaction containing lesser fillerfiller aggregate in the polymer matrix. The study of melt flow rheology is also in line with the DMA strain sweep results. The thermal degradation stability of C60E40/PCL-gHNT(C) nanocomposite is found to be higher than any other nanocomposites, which is because of the homogeneous state of filler dispersion in the matrix with improved polymerfiller interaction. Hence, it can be concluded that in CPE/EMA (60/40) blend system, the in-situ covalent modification of HNT by PCL can act as feasible method for improving polymer-filler interfacial adhesion. Though APTES can also enhance the hydrophobicity of HNT, but PCL being compatible with both CPE and EMA provided an added advantage in this blend system.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Supporting Information Available Table S1: Wide angle X-ray diffraction (WAXD) of pristine and modified HNT. Table S2: Characteristic WAXD peak positions and corresponding d-spacing of Nanocomposites from Figure 7. Table S3: TGA data of neat C60E40 blend and its prepared (nano)composites. This information is available free of charge via the Internet at http://pubs.acs.org/.

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Table Captions Table 1: Sample designations with their composition Table 2: FTIR bands and corresponding assignments of pristine HNT, PCL-g-HNT(C), PCL-g-HNT(P), APTES-g-HNT and PCL

Table 3: Interphase interaction parameter B values (using Equation 6) and effective freespace length Lf* (using Equation 8) of prepared nanocomposites

Table 4: The Glass transition temperature (Tg) and storage modulus E’ values at high temperature (100 °C) of all samples from DMA analysis

Table 1 Sample designations

EMA

HNT

PCL-g-

PCL-g-

APTES-g-

(phr) (phr)

(wt%)

HNT(C)

HNT(P)

HNT

(wt%)

(wt%)

(wt%)

CPE

C60E40/HNT

60

40

5

-

-

-

C60E40/PCL-g-

60

40

-

5

-

-

60

40

-

-

5

-

60

40

-

-

-

5

HNT(C) C60E40/PCL-gHNT(P) C60E40/APTES-gHNT * Other ingredients are quantitatively equal in all samples (MgO, DBTDL, and Irganox 1010 are in 3, 1, and 1 phr respectively) and compositions are in phr w.r.t. polymer



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Page 40 of 51

Table 2 Assignments (cm-1)

HNT

PCL-g-HNT(C)

PCL-g-HNT(P)

APTES-

PCL

g-HNT 3696

3695

3696

3695

-

3618

3618

3618

3618

-

In plane Si-O stretching

1032

1032

1032

1032

-

Deformation of C-H2

-

-

-

1495

-

Deformation of N-H2

-

-

-

1556

-

Perpendicular Si-O

689

685

686

685

-

753

750

752

751

-

-

1724.55

1725.00

1725

1725.43

O-H stretching of innersurface hydroxyl groups O-H stretching of inner hydroxyl groups

stretching Perpendicular Si-O stretching C=O stretching peak


Page 41 of 51

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Table 3

Sample code

B value

Effective free-space length Lf* (nm)

C60E40/HNT

2.34

280

C60E40/PCL-g-HNT(P)

2.44

380

C60E40/APTES-g-HNT

2.68

390

C60E40/PCL-g-HNT(C)

3.36

430

Table 4

Sample

Glass transition

Storage modulus

designations

temperature

E’ at 100 °C

Tg (°C)

(MPa)

C60E40

5.2

27.85

C60E40/HNT

6.0

122.5

C60E40/PCL-gHNT(P)

7.3

460.5

C60E40/APTES-gHNT

9.5

491.9

C60E40/PCL-gHNT(C)

11.5

579.1



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Figure Captions Figure 1: Plausible schematics indicating reaction sites of (a) in-situ ring opening polymerization of (ε-caprolactone) onto HNT surface and (b) covalent grafting of APTES onto HNT surface

Figure 2: (A) XRD patterns of unmodified and modified HNTs: (a) pristine HNT, (b) PCLg-HNT(P), (c) PCL-g-HNT(C), and (d) APTES-g-HNT and (B) FTIR spectra of PCL, modified HNT and pristine HNT (* characteristic peaks of HNT, ** characteristic peaks of PCL)

Figure 3: HRTEM images of (a) pristine HNT in 50 nm scale, (b) nanotube and its open end in 20 nm scale, (c) one open end of HNT in 10 nm scale; (d) PCL-g-HNT(P), (e) PCLg-HNT(C), and (f) APTES-g-HNT

Figure 4: Contact angles of (a) Pristine HNT, (b) PCL-g-HNT(P), (C) APTES-g-HNT, (d) PCL-g-HNT(C)

Figure 5: Column bar diagram of nanocomposites (a) Tensile strength, (b) Modulus at 100% elongation, and (c) Elongation at break of unfilled C6E40 blends and prepared nanocomposites (an average value of 5 samples are taken)

Figure 6: Strain hardening of polymer matrix obtained from the slope of log σT vs. log λ plot

Figure 7: Variation of (a) storage modulus E’ and (b) loss factor tanδ with temperature in neat C60E40 blend and its nanocomposites


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Figure 8: Strain dependence of storage modulus E’ at 25 °C at 10 Hz for neat C60E40 blend and its nanocomposites (* in the inset table E’∝ and (E’0 — E’∝) represents the measure of reinforcement and filler-filler network structure respectively)

Figure 9: Log (apparent shear viscosity, η) vs. Log (shear rate, γ˚) plots of neat C60E40 blend and its nanocomposites

Figure 10: XRD patterns of HNT nanocomposites (a) C60E40/HNT, (b) C60E40/PCL-gHNT(P), (C) C60E40/APTES-g-HNT, (d) C60E40/PCL-g-HNT(C) (* peak of pure HNT disappeared in nanocomposites)

Figure 11: Representative TEM images of (a) C60E40/HNT, (b) C60E40/PCL-g-HNT(P), (c) C60E40/APTES-g-HNT, and (d) C60E40/PCL-g-HNT(C)

Figure 12: FE-SEM images of (a) C60E40/HNT, (b) C60E40/PCL-g-HNT(P), (c) C60E40/APTES-g-HNT, and (d) C60E40/PCL-g-HNT(C) and corresponding TEM images of (e) C60E40/HNT, (f) C60E40/PCL-g-HNT(P), (g) C60E40/APTES-g-HNT, and (h) C60E40/PCL-g-HNT(C) at high magnification

Figure 13: 2D phase images of (a) C60E40/HNT (b) C60E40/PCL-g-HNT(P), (c) C60E40/APTES-g-HNT ,(d) C60E40/PCL-g-HNT(C)

Figure 14: (a) TGA and (b) DTG plots of neat C60E40 blend and their (nano)composites



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Figure 1

Figure 2


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Figure 3

Figure 4



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 5

Figure 6


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Figure 7

Figure 8

Figure 9 47 ACS Paragon Plus Environment


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Figure 10

Figure 11


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Figure 12

Figure 13 49 ACS Paragon Plus Environment


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Figure 14


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