Design of a Molecular Architecture via a Green Route for an Improved

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Design of a Molecular Architecture via a Green Route for Improved Silica Reinforced Nanocomposite using Bio-resources Pranabesh Sahu, Preetom Sarkar, and Anil K. Bhowmick ACS Sustainable Chem. Eng., Just Accepted Manuscript • DOI: 10.1021/ acssuschemeng.8b00383 • Publication Date (Web): 10 Apr 2018 Downloaded from http://pubs.acs.org on April 11, 2018

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ACS Sustainable Chemistry & Engineering

Design of a Molecular Architecture via a Green Route for Improved Silica Reinforced Nanocomposite using Bio-resources

Pranabesh Sahu, Preetom Sarkar and Anil K. Bhowmick*

Rubber Technology Centre, Indian Institute of Technology Kharagpur, Kharagpur -721302, West Bengal, India *Corresponding author: E-mail: [email protected]; [email protected] Tel: +91 3222 283180, Fax: +91 3222 220312

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ABSTRACT Synthesis of sustainable polymers from terpenes via a facile and green approach has evolved as a high-impact research field. In this work, we have designed a new molecular architecture and synthesized a highly reactive epoxy group-functionalized bio-based elastomer, poly (myrcene-co-glycidyl methacrylate), by an environment-friendly emulsion polymerization technique with an ultimate aim to have better reinforcement ability for silica reinforced tire. The copolymers displayed molecular weight in the range of 71,500 to 105,870 Da and a sub-ambient glass transition temperature between −48 to −8 °C. By combining the molecular structure with non-petroleum based silica, silica/elastomer green nanocomposite was designed to investigate the effect of the epoxy groups on the interfacial interaction, morphology, and performance of the nanocomposites. The silica reinforced elastomer vulcanizate exhibited better silica dispersion, higher mechanical properties, and greater traction and wet skid resistance than the pristine elastomer. The current approach provides an effective route to make sustainable elastomers for diverse applications replacing petroleum-based analogues, which would be particularly useful to the automobile industry.

KEYWORDS:

Terpene,

Bio-based

elastomer,

Emulsion

nanocomposite, Mechanical properties.


polymerization,

Silica

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INTRODUCTION Synthesis of polymers from natural resources is driven by immense interest in sustainable development to reduce our dependence on petroleum-based products. Present-day research has focused mainly on replacing fossil raw materials with renewable alternatives to provide an improved quality of life and a cleaner environment.15

With the upsurge of bio-based polymers, bio-based elastomers from natural reservoir

have allured colossal attention.6 Elastomers show vast extensibility up to several hundred percent, and have the ability to recover their original shape and configuration on release of applied stress.7 The application areas of elastomers are numerous and elastomers play a significant role in automobiles, defense and daily living. Among the manifold natural and sustainable elastomeric materials derived from renewable biomass or agricultural waste, terpenes have attracted great interest in current research.8-11 ‘Terpene’ refers to one of the largest families of naturally occurring compounds synthesized by conifers, and various plants.12-14 Among the vast terpene reserves, βmyrcene (MY) has been widely explored in designing elastomers and is a promising building block of synthetic polymers in recent years. It is now considered as a natural base chemical in sustainable chemistry.15 Previously, Sarkar and Bhowmick synthesized an elastomer from β-myrcene via conventional emulsion polymerization and interpreted the microstructure of poly-myrcene in detail.16 Later a series of biobased random rubbery copolymers of β-myrcene with different renewable synthons were also reported.17,18 Recently, synthesis of sustainable methacrylate copolymer series with β-myrcene by emulsion polymerization was also documented.19



Our objective is to use a functional monomer for copolymerization with β-myrcene to develop a multifunctional elastomer matrix which is required for silica reinforcement in next generation tires. Glycidyl methacrylate (GMA) was taken as a comonomer with βmyrcene, as it can provide a wide range of choice for chemical modification of the synthesized polymer. The copolymers based on glycidyl methacrylate belong to the potential category of functional polymers for various applications.20,21 The interest in these classes of copolymers is basically due to the ability of the pendent epoxy group to infiltrate into a wide range of chemical modification reactions.22 Beside synthesis and characterization of functional block copolymers, terpolymer with GMA was reported.23,24 However, till now no research has focused on the synthesis of biobased copolymer of βmyrcene and GMA by a green polymerization method. Our primary aim of the work is to synthesize bio-based elastomer. So, β-myrcene was chosen as the main monomer, as it offers rubbery properties to the polymer. GMA was chosen basically due to its bifunctional nature with a double bond and an epoxy group, to provide functional modification to the polymer and interaction with the silica surface. It was not chosen as the main monomer because a highly crosslinked polymer results with higher amount of GMA. In this work, we mainly focus our attention on copolymerization of β-myrcene with GMA to synthesize a bio-based elastomer poly (MY-co-GMA) (PMG) by redox-initiated emulsion polymerization. Emulsion polymerization was used due to solvent-free, mild reaction conditions and commercially viable green technique.25,26 In addition to evaluation of monomer reactivity ratios, structure-property relationship of the copolymers was established. Here, we also report the application of two-dimensional (2-D) NMR double quantum filtered correlation spectroscopy (COSY and NOESY) to characterize


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the interactions between differentially coupled protons. Thermal properties as well as Xray diffraction study of the copolymers were also investigated. Such studies are important in understanding similar copolymers. Elastomers are generally soft and weak with very poor mechanical properties. The reinforcement of elastomers by finely divided fillers, particularly carbon black and silica, is fundamental to the rubber industry to improve elastomer properties.27 The research on rubber reinforcement is a traditional but vitally essential topic and an important issue in elastomer science and engineering.28-30 Different functionalization of the elastomer matrix was done to observe covalent interaction and dispersion of filler in the matrix.31,32 Silica was used for reinforcement with the synthesized biobased PMG elastomer, as it is epoxy functionalized. Silica was also chosen as it is non-petroleum based and has a wide commercial application in tire industry. It is interesting to note that improvement of silica-elastomer interaction is still a challenging research problem. The silanol groups on the silica surface facilitate covalent bonding interactions with elastomer to promote uniform dispersion of silica. Apart from the detailed synthesis of bio-based functional elastomer,

the

present

study

also

intends

to

prepare

bio-based

elastomer

nanocomposite with improved performance and to study the interfacial interaction between the epoxy group of biobased elastomer matrix and the nano-filler. The ring opening reaction between the epoxy-functionalized elastomer and silica on application of temperature during vulcanization has been also discussed. During the research, it shows that the strong interfacial interaction between the epoxy-functionalized biobased elastomer and silica effectively improves the dynamical and mechanical properties of the nanocomposites than the pristine polymer. 5 ACS Paragon Plus Environment


Material and Methods Raw Materials. β-myrcene (MY, 99%), was purchased from Sigma-Aldrich and purified before use. Glycidyl methacrylate (GMA, 97% purity) was bought from Sigma-Aldrich and passed through a basic alumina column before use. Potassium oleate (K-Oleate, 98%) and Ferric ethylene diamine tetraacetic acid salt (Fe-EDTA) were purchased from Loba Chemie and Alfa Aesar respectively. Potassium phosphate tribasic (K3PO4), potassium chloride (KCl), sodium hydroxymethanesulfinate (SHS) and tert-butyl hydroperoxide (Luperox® TBH70X) were procured from Sigma-Aldrich and used as received. Deionized water (DI H2O) was used for all the experiments. Redox initiated emulsion copolymerization of poly (MY-co-GMA) (PMG). Bio-based PMG elastomer was synthesized via redox-initiated emulsion polymerization according to the recipe in Table 1. DI water, potassium oleate, potassium chloride and potassium phosphate tribasic buffer were charged into a round bottom flask and stirred at 300 rpm for 20 min. Subsequently, β-myrcene and glycidyl methacrylate in appropriate amount were added into the RB flask and allowed to mix for further 20 min. Fe-EDTA and SHS were injected into the reaction mixture upon attaining a stable emulsion. Thereafter, the reactor was vacuumed and flushed with nitrogen to make an inert atmosphere, followed by the addition of tert-butyl hydroperoxide (TBHP) solution into the flask. The polymerization was allowed to proceed for 12 h at 20 °C. The obtained latex was coagulated using excess ethanol. The polymer then was washed thoroughly with DI water and dried in a vacuum oven for 24 h at 45 °C to obtain a rubbery type polymer. A few series of copolymers with varying contents of MY and GMA were also prepared following the same recipe. The physical appearance of the copolymers is shown in 6 ACS Paragon Plus Environment

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Figure S1 (Supporting information, SI). The polymerization mechanism is shown in Scheme 1. The kinetic study was done by taking aliquots from the reaction mixture at various time intervals within 4 h and the reaction rate was calculated. Table 1: Recipe for Redox-Initiated Emulsion Polymerization. Ingredients β-myrcene glycidyl methacrylate DI water potassium oleate potassium chloride (KCl) potassium phosphate tribasic (K3PO4) Fe-EDTA sodium hydroxymethane sulfinate(SHS) tert-butyl hydroperoxide (TBHP)

Amount (g, in phra) Variable Variable 180 4.5 0.3 2.0 0.15 0.05 0.06

a

phr : parts per hundred parts of rubber

Scheme 1. Redox Emulsion Polymerization of poly (MY-co-GMA). 7 ACS Paragon Plus Environment


Preparation of PMG/Silica Green Nanocomposites.

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Among

all the synthesized

copolymers, poly (MY80GMA20) (designated as PMG-20 later) was chosen initially for compounding purpose, where the subscript denotes the weight percent of the respective co-monomer. Standard sulfur cure system was used to vulcanize the synthesized elastomers. AEROSIL-200 having specific surface area of 200±25 m2/g was used for reinforcement

purpose.

The

compounding

formulation

for

the

preparation

of

nanocomposite is given in Table 2. Gum poly-myrcene elastomer and poly (MY90GMA10) (designated as PMG-10) were also vulcanized and tested using the same formulation for comparison purpose. At first, the synthesized elastomers were mixed with silica using a 6 inch/12 inch two-roll mill (Schwabenthan, Berlin) at 25 °C. Then, all the compounding ingredients were mixed sequentially to obtain the rubber nanocomposites. The optimum cure time of the compounded sample was determined by a Monsanto Rheometer R100S machine, USA at 150°C. Finally, the compound was vulcanized in a hydraulic compression molding machine (David Bridge Company, England) at 150 °C under 15 MPa pressure for its optimum cure time as obtained from an oscillating disc rheometer. Table 2: Compounding Formulation for Silica Nanocomposites. Ingredients

Loading (phr)a

polymer silica (AEROSIL-200) zinc Oxide stearic Acid i-PPDb accelerator TBBSc sulfur

100 20 5 2.0 1.0 1.2 1.5

a

parts per hundred parts of rubber.

b

N-2-Propyl-N'-phenyl-p-phenylenediamine, c N-tert-butyl-2-benzothiazole sulfenamide. 8 ACS Paragon Plus Environment

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Instrumentation. The number average molecular weight (Mn) of the polymers was measured by gel permeation chromatography (GPC) at 25 °C using an SHIMADZU-GPC instrument, having a refractive index detector equipped with Phenogel 10u Columns and using Tetrahydrofuran (THF) as the eluent (sample concentration 1 mg/mL) at a flow rate of 1 mL/min. Polystyrene standard was used for calibration. The gel percentage of the copolymers was calculated after extraction using THF for 20 h at room temperature (25 °C).The value was taken as the ratio of dried polymer weight to its original value. The particle size and its distribution was measured by a dynamic light scattering (DLS) method using a Malvern Nano ZS instrument employing a 4 mW He−Ne laser (λ = 632.8 nm) at a scattering angle of 90°. The Fourier transform infrared (FT-IR) spectra of the monomers, synthesized homopolymers and copolymers were recorded in a PerkinElmer Spectrum 400 machine (resolution 4 cm−1) within a spectral range of 4000−400 cm−1 using an attenuated total reflectance (ATR) sampling technique. 1

H and 13C nuclear magnetic resonance (NMR) spectra were recorded on an AVANCE III

400 Ascend Bruker instrument operating at 600 MHz. Two-dimensional (2-D) NMR COSY and NOESY measurements were made to collect two-dimensional hyper-complex data to assign different proton-proton spin-spin coupling and structural information of the polymer. After weighing with a shifted sine-bell function, the data were Fourier transformed in the absolute value mode. Chloroform-D (CDCl3) was used as the solvent and the chemical shift values were reported in δ (ppm).



Differential scanning calorimetry (DSC) measurements were recorded using a NETZSCH DSC 200F3Maia instrument at a heating rate of 10°C/min under N2 atmosphere. The samples were heated from −100 to +100 °C. Glass transition temperatures (Tg) were determined from the second heating run, taken as the mean value between the onset and end-point temperatures. The thermal stability of the polymers was determined by thermogravimetric analysis (TGA), conducted on a SDT Q600 TA instrument. The samples were scanned in the temperature range of 30–800°C at a heating rate of 10 °C/min under N2 atmosphere. Wide angle X-ray diffraction (WAXD) analysis was carried out using Nickel filtered copper Kα radiation source (0.154 nm) at room temperature in the 2θ range of 10°−50° with a scan rate of 3°/min using a 2nd generation D2 PHASER, BRUKER instrument. Characterization of the Silica Nanocomposite. The tensile tests of the silica-filled nanocomposites were performed according to ASTM D 412 standard using a Zwick/Roell Z010 universal testing machine at room temperature. Dumbbell-shaped specimens were cut from the moulded sheets and the tensile strength was measured at a crosshead speed of 20 mm/min. The length between the two pneumatic grips was 20 mm. The data were analyzed by testXpert II software of the Zwick/Roell instrument and average of these measurements was taken as the final result. Dynamic mechanical properties of the vulcanizates were determined using a METRAVIB 50N (France) dynamic mechanical thermal analyzer in tension mode. Temperature sweep experiments were carried out at 1 Hz frequency and 0.1% dynamic strain. The scanned temperature range was from −100 to +100 °C at a ramp rate of 3 °C/min. 10 ACS Paragon Plus Environment

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The swelling property of the cured samples was measured by equilibrium solvent swelling experiment33 in a suitable solvent. A specified test sample of particular dimension, weighed Wo, was immersed into Tetrahydrofuran (THF) solvent. The sample was kept at room temperature till swelling equilibrium was reached (when the mass remained almost constant). After reaching equilibrium (about 72 h) the sample was taken out, mopped with filter paper to remove the excess solvent, and weighed (W 1). Thereafter, the sample was dried in an oven at 60 °C for 24 h to get the deswollen weight (W 2) of the sample. The swelling ratio Q was calculated by

where ρs and ρr are the densities of the solvent and the elastomer respectively. The density of the elastomer (ρr) was measured by using Wallace High Precision Densimeter following the ASTM D297-93-16 (Rubber) standard. The morphology and dispersion of silica nanoparticles was observed for cryofractured sample using ZEISS Field Emission Scanning Electron Microscopy (FESEM) at an acceleration voltage of 5 kV. Gold coating was done on fractured surfaces before the experiment.



RESULTS AND DISCUSSION Synthesis of redox initiated poly (MY-co-GMA) (PMG) copolymer. The bio-based PMG elastomers were successfully synthesized by free radical redox emulsion polymerization. PMG copolymers were prepared with different monomer ratios at 20 °C under atmospheric pressure for 12 h. Low polymerization temperature gives the advantage of lower chain branching and crosslinking in the synthesis of polymers. Higher reaction temperature yields highly crosslinked polymer due to the presence of epoxy group in GMA. So, low temperature polymerization process was optimized by varying the reaction temperature and keeping the ingredients same. K-Oleate was used as surfactant, K3PO4 as a buffer, sodium hydroxymethanesulfinate (SHS) as an oxidant, FeEDTA as reductant and tert-butyl hydroperoxide (TBHP) as a redox initiator. The oxidant, SHS, and the reductant, Fe-EDTA, were used to lower the activation energy of TBHP radical decomposition. The molecular weights of all the obtained polymers were confirmed at the end of the polymerization. The molecular weight of the copolymers was determined sequentially by GPC measurement. Optimization of Conditions. In order to optimize the reaction conditions, the effect of reaction time on molecular weight and yield of the copolymer was measured at 4 h time interval up to 16 h reaction time. Figure 1a represents the variation of yield percentage and molecular weight with time for a fixed MY /GMA mass ratio of 90/10 copolymer. Reaction temperature was kept fixed (at 20 °C) to avoid crosslinking between the polymer chains due to the presence of highly cross-linkable epoxy group in GMA. With increasing reaction time, the yield increases reaching a maximum of 90% at 12 h and the molecular weight increases from 4 h to 12 h, thereafter showing a decreasing trend. This 12 ACS Paragon Plus Environment

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decrease can be attributed to the scission of the macromolecular chains at higher reaction time resulting in the formation of oligomers. Accordingly, the reaction condition was set to 12 h reaction time at 20 °C and was used for the synthesis of the homopolymers and copolymers.

A series of copolymers having different weight

percentages of the β-myrcene moiety were synthesized. Physical appearance of the representative samples is shown in Figure S1 (SI). The yield percentage, molecular weight, gel content, PDI and latex particle size of the synthesized copolymers are presented in Table 3, where the subscript denotes the weight percent of the respective co-monomer. The yield of the polymers obtained is in the range from 65% to 84%. The gel content of the copolymers increases with increase in GMA content due to the formation of crosslinks between the macromolecular chains and epoxy ring of the GMA unit. With the increase of GMA weight percent, the number average molecular weight (Mn) of the copolymer decreases when these are compared with that of individual homopolymers, poly (MY100GMA0) and poly (MY0GMA100) respectively. It is evident from Table 3 that among the series of copolymers, poly (MY90GMA10) showed highest molecular weight of 105,870 Da. With increase in GMA content, the molecular weight decreased up to 71,500 Da for poly (MY70GMA30) copolymer and then shows an increment to a maximum of 94,040 Da for poly (MY50GMA50). Beyond this concentration of GMA, the copolymer formed gel and were difficult to process. The polydispersity index (PDI) also decreases with decrease in β-myrcene content in copolymer series. The Z-average particle size of the poly (MY-co-GMA) latices was obtained in the range of 92-249 nm.



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The variation in particle size with the copolymer composition was studied by monitoring the reaction at certain intervals of time (2h to 12 h) for poly (MY90GMA10). With increase in time, the particle size increases and the copolymer composition varies from 83/17 to 88/12 MY:GMA ratio (Table S1, SI).

Figure 1. (a) Time dependence of poly (MY90GMA10) synthesis on yield and molecular weight (b) Kinetic plots of PMY, PGMA and poly (MY90GMA10) copolymer. Table 3. Molecular Weight, Yield, Gel Content, PDI, and Particle Size of Redox Initiated poly (MY-co-GMA) Copolymers.

Polymer

Yield Gel Content (%) (%)

Molecular Weight (Da)

Mw/Mn (PDI)

ZAverage diameter (nm)

poly (MY100GMA0)

65

5

168,250

3.41

92

poly (MY90GMA10)

78

11

105,870

4.87

192

poly (MY80GMA20)

82

14

72,400

2.05

135

poly (MY70GMA30)

76

17

71,500

2.05

157

poly (MY50GMA50)

84

18

94,040

1.96

249

poly (MY0GMA100)

80

20

160,140

1.47

230


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Kinetic curves were recorded by taking aliquots from the reaction mixture at various time intervals and then determining the monomer conversion. The kinetic plot in Figure 1b shows that the dependence of ln (1/1-X) {X = conversion of monomers, X=obtained yield of the polymer (g) /amount of monomer (g) in the aliquot taken each time} versus time is linear for poly-myrcene (PMY), poly-GMA (PGMA) and poly (MY90GMA10) copolymer. The reaction rate was determined from the slope of the conversion-time curve that corresponded to the pseudo steady-state rate of polymerization, as described by the Smith-Ewart theory of emulsion polymerization.34 It is apparent from Figure 1b that the rate for homopolymerization of GMA is much higher (k =2.70 × 10−2 min−1) compared to the rate of homopolymerization of β-myrcene (k = 2.37 × 10−3 min−1) because of its solubility in water, whereas the rate for copolymerization (k =1.00 × 10−2 min−1) lies in between the two. This is explained with the help of micellar nucleation mechanism34,35 and relative solubility of the monomers in the aqueous phase. Figure 2 represents a schematic representation of various plausible intervals of the emulsion copolymerization reaction. The addition of GMA has an incremental effect on the rate of copolymerization (compared to the homo-polymerization of β-myrcene). As the GMA monomer was introduced into the polymerization system [poly (MY90GMA10)], the molecular weight of the resulting copolymers was found to decrease initially compared to that of the individual homopolymers. Then, it exhibits an increase up to a value of 94,040 Da for poly (MY50GMA50).

Being a unique and simple chemical process, emulsion

copolymerization is a combination of several complicated mechanistic events. Several parameters including polarity, monomer partitioning, and water solubility of the monomers are the key factors that govern the particle growth stage in emulsion



copolymerization. The reaction medium consists of the following systems: MY and GMA monomer reservoirs, monomer swollen micelles, surfactant molecules, and water-soluble initiator. Due to polar nature, GMA monomer has relatively higher water solubility (24.64 vs 16.69 MPa1/2 for MY at 25 °C) as determined from the theoretical solubility parameter calculation36 (see Supporting Information, Table S2). So, GMA has relatively higher tendency to be miscible in water. The waterborne free radicals after the decomposition of the initiator first react with the more water-soluble GMA monomer dissolved in the continuous aqueous phase and produce the oligomeric radicals (Interval I). This would result in the increased hydrophobicity of oligomeric radicals after a critical chain length is achieved and tend to enter monomer-swollen micelles to continue propagation process. However, at the same time the polar oligomers will be less prone to penetrate into the micelles and the presence of electron withdrawing carbonyl groups make the GMA intermediate radical more stable than that of the MY, thereby slowing down the chain propagation. This phenomenon obstructs the growth of the polymer chain and thus also reduces the molecular weight of the copolymer. It was believed that due to hydrocarbon nature of myrcene, the interiors of most of these micelles were predominantly occupied by MY monomers. As a consequence, monomer-swollen micelles are successfully transformed into particle nuclei (Interval II). So, the penetration of the MY monomer molecules would be facilitated and hence a higher growth of the polymer chains and rate of polymerization would be observed. For 90/10 MY/GMA copolymer, most of the GMA units are spent in making the active radicals at the nucleation stage. Due to hydrocarbon nature, the MY monomers would tend to form larger monomer reservoirs by self-association. This leads to an increased 16 ACS Paragon Plus Environment

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set of active radical formation during the nucleation step. Since MY constitutes the major share in the copolymer, it influences the chain growth significantly along with the presence of GMA unit. The concentration of monomer in the reaction loci continues to decrease, resulting in the consumption of monomer droplets from monomer reservoirs. So, the reaction loci become monomer starved towards the end of polymerization and the polymerization rate decreases (Interval III).

Figure 2. A schematic representation of emulsion copolymerization of MY and GMA. Reactivity Ratios of β-myrcene and Glycidyl methacrylate. The reactivity ratios of βmyrcene and glycidyl methacrylate in the redox initiated emulsion polymerization were determined by two linear methods, mainly the Fineman−Ross (FR)37 method and Kelen−Tüdös (KT)38 method. Reactivity ratios were calculated by carrying out 17 ACS Paragon Plus Environment


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polymerization at varying monomer ratios (approximately 10−15% conversion) so that the changes in the monomer feed composition are minimal. The FR and KT plots are shown in Figure 3. The detailed theoretical calculation is discussed in supporting information and the relevant parameters for the copolymers are presented in Table S3 (SI), which summarize the reactivity ratios as obtained by both the methods. Although the trend in reactivity ratio values obtained from the two methods is quite similar, it is evident that the reactivity ratio of MY is greater than that of GMA. This suggests that the MY monomer is more reactive than GMA monomer towards both the propagating species (MY* and GMA*). But for both the methods,

and

are less than one. So,

the polymerization of MY and GMA belongs to an azeotropic copolymerization with an average azeotropic point 0.6 (Table S4, SI). The copolymerization, if carried out at a MY molar fraction of 0.6, will produce a copolymer of the same composition as the feed. The set of conditions arises due to copolymerization of monomers of different polarity. Due to large difference in polarity and electron densities (C=C double bond) between two monomers, the product

×

decreases.39,40

Figure 3. Reactivity ratios of MY and GMA in redox emulsion polymerization. (a) Fineman-Ross method and (b) Kelen-Tüdös method. 18 ACS Paragon Plus Environment

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FTIR Spectra of monomers, homopolymers and PMG copolymers: Figure 4a presents the FTIR spectra of GMA, MY monomers and their synthesized homopolymers, PGMA and PMY respectively. The absorption peaks at 3000, 2958 and 2932 cm-1 in the GMA monomer are due to =C-H, -CH3 and –CH2 stretching vibrations respectively. The peak at 3000 cm-1 (=C-H stretching vibration) disappears upon polymerization. The peaks at 1716 cm−1 and 1638 cm−1 belong to the C=O and C=C stretching vibrations of GMA. The disappearance of C=C peak after homopolymerization depicts the consumption of double bond confirming the formation of polymer, which can be observed in the PGMA spectra. The peak due to C=O stretching is preserved and appears at a higher value of 1726 cm-1 in the PGMA polymer. The peak at 907 cm-1 corresponding to ring vibration of the epoxy group is also preserved after the polymerization, indicating the presence of unaffected epoxy group. The peak at 842 cm-1 is attributed to =C-H bending vibration of the double bond in GMA. The weak peak at 3090 cm-1 in the MY monomer represents the =C-H stretching vibration which disappears on polymerization. The peaks at 2968, 2920, 2858 cm-1 in MY monomer are due to –CH3, -CH2, -CH asymmetric stretching vibrations respectively which are broadened in PMY spectra. The peak at 1642 cm-1 is due to isolated double bond (C=C stretching vibration) present and the strong peak at 1595 cm-1 is assigned to conjugated double bond (C=C stretching) present in the monomer. The complete disappearance of 1595 cm-1 band indicates the participation of conjugated double bond in the polymer formation and the retention of 1642 cm-1 peak indicates that the isolated double bond is intact during the polymerization. The absorption peaks at 1445 and 1376 cm-1 in the MY monomer and the PMY polymer are attributed to –CH2 and –CH3 bending vibrations respectively. The



strong absorption peaks around 990 and 891 cm-1 in MY are due to sp2 C–H bending vibration of conjugated double bonded carbon centres respectively. The sharp decrease in the intensities of these peaks in poly-myrcene indicates the consumption of the conjugated double bond in the polymerization process. The peak at 827 cm-1 is attributed to sp2 C–H bending vibration of the isolated double bond carbon which remains unchanged even after polymerization. In the poly (MY90GMA10) copolymer spectra, the double bond stretching absorptions of GMA and MY monomer at 1638 and 1595 cm-1 disappeared after the polymerization, confirming that the monomers (GMA and MY) were successfully copolymerized. The presence of small peak at 1655 cm-1 was because of the isolated double bond present in MY monomer which did not take part in the copolymerization reaction. The peaks due to C=O stretching, ring vibration of epoxy group in GMA at 1730 and 907 cm-1 are preserved even after copolymerization indicating the incorporation of GMA in the copolymer. The dual absorption peaks at 1448 and 1376 cm-1 are due to –CH2 and –CH3 bending vibrations of MY, which are well preserved in the copolymer also. The broad peaks at 992 and 840 cm-1 are attributed to sp2 C-H bending peak present in the copolymer. The broad bands at 2855, 2921 and 2964 cm-1 are attributed to the asymmetric stretching vibrations of the –CH,-CH2 and –CH3 group of the synthesized poly (MY-co-GMA) polymer. Figure 4b shows the absorption spectra of the synthesized poly (MY-co-GMA) copolymer series with varying GMA ratio. It is evident that the peak intensity at 1730 cm−1 (C=O stretching of ester bond) and 907 cm-1 (ring vibration of epoxy group) increases with increase in the GMA content. 20 ACS Paragon Plus Environment

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Figure 4. FTIR spectra of (a) Homopolymers and copolymers. (b) poly (MY-co-GMA) copolymers with varying GMA ratio. High-resolution NMR spectroscopy of homopolymers and poly (MY-co-GMA) Copolymer. The structural characterization of the homopolymers and the copolymers was

performed using NMR spectroscopy. A set of both one-dimensional (1D) and two-

dimensional (2D) experiments was conducted to assign the spectra of the synthesized homopolymers and poly (MY-co-GMA) copolymer. The 1H NMR and

13

C NMR spectra of

the monomers and their homopolymers were also taken for interpretation of copolymer microstructure and the detailed assignments of each peak to the molecular structure are shown in the supporting information (Figure S2−S4). The NMR spectrum of the characteristic copolymer, poly (MY90GMA10) was also taken (Figure 5). NMR spectra (1H NMR and

13

C NMR) of another copolymer poly (MY80GMA20) are also shown in the

supporting information (Figure S5). The spectral signals are well assigned to various magnetically different protons and carbons. The chemical shift values of the relevant peaks in 1H and 13C NMR spectra are listed below.



1

H NMR Data of poly (MY90GMA10) (CDCl3, 600 MHz): δ-5.13 (4H, 3, 7, 3′ and 7′- H),

4.38 & 3.88 (2H, d-H), 3.20 (1H, e-H), 2.83 & 2.64 (2H, f-H), 1.69 & 1.61 (12H, 9, 10, 9′,and 10’-H), 2.05 (16H, 1, 4, 5, 6, 1’, 4’, 5′,and 6′-H),1.28 (2H, a-H) and 0.90 (3H, b’-H). 13

C NMR Data of poly (MY90GMA10) (CDCl3, 600 MHz): δ-176.8 (g-C), 131.5 (2,2’,8 and

8’-C), 124.0 (3,7,3’ and 7’−C), 65.0 (d-C), 49.0 (e-C), 46.1 (b-C), 44.5 (f-C), 37.3 (5 and 5′− C), 31.0 (1 and 1′-C), 26.9-25.6 (4,4’,6 and 6’-C), 23.1-20.5 (a,9,9’,10 and10’-C),17.6 (b’-C). In the 1H NMR spectra of GMA (Figure S2a, SI), the peaks at δ = 6.12 and 5.57 ppm represent the two olefinic hydrogens, which disappear completely after polymerization, indicating the consumption of the double bond during PGMA formation (Figure S4a, SI). The two distinct doublet peaks at 4.45 and 3.96 ppm arise due to magnetically nonequivalent methylene group protons adjacent to epoxy ring, which remain unaffected after polymerization. The small peak at 3.22 ppm appears as multiplet and the two distinct doublet peaks at 2.83 and 2.64 ppm originate from the protons of epoxy group which remain unaltered, indicating that the epoxy ring is preserved even after polymerization. From the13C NMR spectra of GMA (Figure S2b, SI), the peaks at δ = 135.7 and 126.0 ppm correspond to the double bonded carbon centers, which vanish after polymerization (Figure S4b, SI). The signature peak at 166.8 ppm is due to carbonyl carbon of ester and the signals at 49.1, 44.4 and 65.1 ppm arise from the epoxy ring carbon and its adjacent methylene carbon, all of which remain unaffected and appear slightly at downfield region after polymerization. According to the previous NMR studies of poly-myrcene16, the spectral assignments of synthesized PMY in this study with different redox recipe are in good record with the available literature. As shown in 22 ACS Paragon Plus Environment

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Figure S4c (SI), the 1H NMR spectrum of redox initiated PMY shows a single peak at 5.14 ppm corresponding to olefinic protons (=CH–units), confirming the participation of conjugated double bond in monomer for polymerization. The methylene protons appear as a broad peak at 2.05 ppm and the methyl protons attached to C-8 in poly-myrcene appear at slight upfield values relative to the monomer (Figure S3, SI) due to the generation of long chain macromolecules. Figure 5a shows the 1H NMR spectra of the copolymer, poly (MY90GMA10) at a feed ratio of 90/10. The broad peak at 5.13 ppm corresponds to the olefinic protons (=CH–units) of poly-myrcene moiety in the copolymer. After copolymerization, the signals for the two olefinic protons at δ = 6.12 and 5.57 ppm of GMA monomer disappear completely, indicating the consumption of the double bond for copolymer formation. The signature peaks at 3.20, 2.83, 2.64, 4.38, 3.88 ppm due to the epoxy ring protons and its adjacent methylene protons indicate the incorporation of GMA unit into the copolymer. All the methylene protons of both MY and GMA moiety appear as usual as in the respective homopolymer. Figure 5b shows the

13

C NMR spectrum of poly (MY90GMA10) copolymer.

The peaks at δ = 131.5 and 124.0 ppm corresponds to the double bonded carbon (polymyrcene moiety) in the copolymer. The peaks for two olefinic carbons at δ = 135.7 and 126.0 ppm of GMA monomer fade away completely. The presence of distinct signal for C=O carbon at δ =176.8 ppm and epoxy group (δ =49.0, 44.5 ppm) and adjacent methylene carbon (δ =65.0 ppm) points to the successful incorporation of the GMA component into the copolymer. The chemical shift values from δ = 17.6 to 37.3 ppm include

13

C signals due to methyl and methylene carbons of both β-myrcene and glycidyl



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methacrylate moieties in the copolymer. The results of the NMR characterization verify that MY and GMA have been successfully copolymerized.

Figure 5. (a) 1H NMR spectrum of the poly (MY90GMA10). (b) poly (MY90GMA10).


13

C NMR spectrum of the

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2-D NMR Characterization. The idea of two-dimensional NMR was introduced to correlate the spectra of coupled nuclear spins. It has proven to be one of the most informative techniques used for the investigation of polymer structure. Among the twodimensional experiments, COSY and NOESY (homonuclear

1

H−1H long-range

correlation) spectra were mainly interpreted in this work. Both F1 and F2 domains of the spectrum are governed by proton resonance frequencies. To interpret the structure of the copolymer, 2-D spectra for individual homopolymers (Figure S6, SI) were first carried out. Figure S6a (SI) represent the COSY-NMR spectra of PGMA, showing that the diagonal peaks have the same frequency coordinate on each axis and appear along the diagonal of the plot, while cross-peaks have different values and appear off the diagonal. The cross peaks determine the couplings between the two nuclei. We found that the epoxy ring protons (e and f-H) are coupled to each other confirming the magnetization transfer between the adjacent protons through chemical bond. The methylene (-CH2) protons signal (d-H) was coupled with the adjacent epoxy ring proton (e-H).The geminal protons (d-H) of methylene group coupled together, giving two symmetrical cross-peaks above and below the diagonal. Figure S6b (SI) presents the COSY spectrum of PMY showing both diagonal and cross-peaks. We found that the olefinic signal (1H each) of 3, 7, 3’, 7’-H were coupled with their adjacent methylene hydrogen present, showing correlation spot respective to the diagonal. Finally, the COSY spectrum (Figure 6) of the copolymer, poly (MY90GMA10) depicts that it is just the combination of two homopolymer correlation spot showing literally all the symmetrical cross-peaks above and below the diagonal. The additional coupling between the protons (4 and 4’-H) of MY moiety with the protons (a-H) of GMA moiety suggest the successful formation of the copolymer.



Figure 6. COSY (1H-1H) spectrum of poly (MY90GMA10) copolymer. Two-dimensional NOESY spectra (1H-1H long range correlation) were acquired in the phase sensitive mode for both homopolymers and copolymers to provide information about the spatial coupling between magnetically different protons that are in close proximity to each other in the polymer. The cross peaks appear off the diagonal indicates couplings between the pairs of nuclei that are spatially close to each other. As can be seen in Figure S7a (SI), the epoxy ring protons (e and f-H) and the adjacent methylene protons (d-H) are spatially coupled with each other giving high-intensity cross peaks. We also found that the methylene (a-H) and methyl (b’-H) protons are also spatially coupled with the epoxy ring protons (e and f-H) and the adjacent methylene protons (d-H) showing correlation spot above and below the diagonal. Figure S7b (SI) shows the NOESY spectrum of PMY with cross-peaks connecting resonances from nuclei that are spatially close to each other. The spatial coupling between the olefinic protons (3, 7, 3’, 7’-H) with the magnetically different methylene and methyl protons suggests that the 26 ACS Paragon Plus Environment

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nuclei are spatially close to each other in the polymer. The correlation spot between the protons of the methyl group (9 and 10-H) could not be well-assigned due to overlapping with the diagonal peak. Figure S8 (SI) depicts the NOESY spectrum of poly (MY90GMA10) copolymer that shows the cross peaks obtained was in good agreement with the homopolymer correlation spot discussed earlier. The spatial coupling between the protons (4 and 4’-H) of MY moiety with the protons (a-H) of GMA moiety was not well observed due to the overlapped diagonal peaks. But the presence of high intensity dispersed cross peaks in low frequencies justifies significant spatial interaction between the methylene protons (4, 4’ and a-H) of both units connecting the copolymer chain. Thermal Properties of Homopolymers and poly (MY-co-GMA) (PMG) Copolymers. For a polymer, glass transition temperature (Tg) plays an important role to identify whether it is rubbery or crystalline. To get the rubbery property Tg should be below theambient temperature. The Tg value of the homopolymers (PMY and PGMA) and the copolymers (with varying GMA content) was determined by DSC. The DSC run from −100 to +100 °C for all the polymer samples indicates a completely amorphous nature of the synthesized materials. Figure 7a shows the DSC thermogram of the pristine polymers. The glass transition temperature of PGMA and PMY was observed at +72 °C and -70 °C respectively. In Figure 7b, the Tg values of the copolymers series lie in between −48 and −8 °C. With increasing GMA content in the copolymer composition, the Tg value of the copolymers increases indicating a decrease in polymer chain flexibility and increase in stiffness of the polymer. The Tg value for the poly (MY90GMA10) shows sufficiently low Tg and is observed at −48 °C. 27 ACS Paragon Plus Environment


Figure 7. DSC thermogram of (a) PMY and PGMA and (b) poly (MY-co-GMA) copolymers. The degradation properties of the synthesized polymers were evaluated using thermogravimetric analysis (TGA). Figure S9 (Supporting Information) represents the thermal decomposition curves for the virgin polymers and their copolymers. Among the pristine polymers (inset of the figure), PGMA displays a single degradation peak and the redox initiated poly-myrcene (PMY) shows two step degradation pattern which is in good agreement with the previous study as well.16 In the case of PMY and PGMA, the sample weight gradually falls after onset of degradation (temperature corresponding to 5% weight loss). After the start of degradation, weight of the PGMA and PMY sample falls gradually, reaching maximum (Tmax) at ∼ 300 °C for PGMA polymer and the second stage of degradation of PMY polymer involves major degradation which starts at 360 °C and gets completed at ∼ 450 °C respectively. The synthesized copolymers exhibit a primary increment in the onset of degradation temperature (T5 ∼ 200 °C) compared to the homopolymers due to the incorporation of glycidyl methacrylate unit and then 28 ACS Paragon Plus Environment

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gradually the copolymer sample weight decreases reaching maximum at ∼ 400 °C. The Tmax values of copolymers are less than the Tmax of PMY due to the possible formation of cross-linked structure by PMY upon degradation. Wide-Angle X-ray Diffraction Analysis. From Figure S10 (SI) it is clear that each copolymer composition shows a broad hump around 2θ = 18°, which mimic the diffraction pattern of natural rubber as reported in literature.41 So, all the synthesized copolymers show humps of low intensity corresponding to their amorphous regions, implying amorphous property. The presence of broad hump of high intensity confirms the amorphous nature of the PGMA polymer. Thus, it is confirmed that all of the synthesized copolymers are amorphous in nature. The presence of pendent alkyl side chain makes the copolymers as well as the homopolymers amorphous. Characterization of PMY/Silica and PMG/Silica Green Nanocomposites. To investigate the reaction of silica with the epoxy group of PMG polymer, the FTIR spectra was recorded after the preparation of the nanocomposite. The corresponding spectra with the epoxy ring region are shown in Figure S11 (SI). It is clear from the figure that the area of the epoxy group vibration peak at 908 cm-1 decreased in the case of PMG20/Silica compound compared to neat PMG-20 polymer. So, after the reaction with silica the intensity of the epoxy ring vibration peak continuously decreases and the occurrence of O-H stretching peak (3428 cm-1) indicates the ring opening reaction between the epoxy ring of polymer and the hydroxyl group of silica during the vulcanization process and interaction (shown in Scheme 2). So, the results of the corresponding spectra suggest that the reaction between silica and PMG polymer occurs during the vulcanization processes.42 29 ACS Paragon Plus Environment


Scheme 2. Ring opening reaction between polymer chains and silica during vulcanization. Curing Characteristics and Mechanical Properties of Nanocomposites. Elastomers are generally soft viscous material, so curing is necessary to convert it into a hard useful engineering product. So, to monitor the curing progress and processing characteristics of elastomer vulcanizate, a typical rheometer curve was obtained from ODR instrument. Figure 8a shows the vulcanization curves of the unfilled and silica filled elastomer nanocomposites at 150 °C. The curing parameters are summarized in Table 4.


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Figure 8. (a) Curing curves and (b) Stress-strain curves of PMY, PMG-20, PMY/Silica and PMG-20/Silica nanocomposites

Table 4. Curing parameters of PMY, PMG-20, PMY/Silica and PMG-20/Silica nanocomposites. Sample

Optimum cure time (min)

Scorch time (min)

ML (dN.m)

MH (dN.m)

PMY

11.15

7.07

2.87

26.53

PMG-20

44.63

0.96

3.09

37.21

PMY/Silica

4.21

2.27

2.54

35.01

PMG-20/Silica

43.27

2.78

3.81

64.01

The minimum torque (ML) of the unfilled PMY and PMG-20 are nearly the same. However, the optimum curing time of unfilled PMG-20 is longer than that of PMY and the maximum torque (MH) is also higher. As all the rubber ingredients and processing are same, the increase in curing time is due to the presence of the epoxy groups in the



PMG-20 polymer, which may react with the rubber accelerators. So, the accelerating effect on vulcanization gets reduced.43-44 The minimum torque (ML) of the PMG-20/Silica compound obtained was higher than PMY/Silica compound, which is due to the reaction between epoxy groups present in the PMG polymer chain and the hydroxyl groups of silica during the mixing. The incorporation of silica extends the curing time of PMG20/Silica compared to PMY/Silica compound due to the possible interfacial interaction of adsorbed rubber accelerators on the silica surfaces with the epoxy groups of PMG polymer. PMG-20/Silica compound exhibited marching cure behavior due to continuous increase of torque with vulcanization time, whereas PMY/Silica shows reversion with cure time. The increase in torque for GMA filled silica composite with respect to pristine PMY silica composite can also be attributed to better dispersion of silica and larger interfacial interaction between the PMG polymer and silica. The mechanical properties of the silica nanocomposites are summarized in Table 5. The representative stress-strain curve of PMY, PMG-20, PMY/Silica and PMG-20/Silica are displayed in Figure 8b. Although the tensile strength of the unfilled PMG-20 was slightly higher, the Young’s modulus and modulus at 100 % strain increased significantly on GMA incorporation. The addition of silica improved the mechanical properties of unfilled polymer vulcanizate further. PMG-20/Silica exhibited better performance than PMY/Silica due to the improved dispersion of silica particle and interfacial interaction with the elastomer matrix. So, the tensile strength was higher for PMG-20/Silica polymer nanocomposite. The addition of silica in PMG-20 elastomer resulted in a strong covalent interaction with the filler-elastomer network, and hence an increase in modulus at 100% strain compared to PMY/Silica compound. Due to high reinforcing efficiency of silica in 32 ACS Paragon Plus Environment

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PMG-20 polymer, the slippage of polymer chains was significantly restricted. So, the elongation at break was lower and the modulus was higher for PMG-20/Silica nanocomposite. An improvement of rigidity of PMG-20/Silica nanocomposite, led to a decrease of ductility and elongation at break of the elastomer vulcanizate. However, the tensile strength of PMG-20/Silica was lower compared to other conventional rubber composites.45-47 Table 5. Mechanical Properties of PMY, PMG-20, PMY/Silica and PMG-20/Silica nanocomposites. Sample

Tensile

Elongation at

Modulus at

Young’s

Strength

break

100%Strain

modulus

(MPa)*

(%)*

(MPa)*

(MPa)*

PMY

0.75±0.08

175±26

0.43±0.04

0.40±0.03

PMG-20

1.07±0.26

107±12

1.02±0.02

0.97±0.18

PMY/Silica

0.93±0.14

160±35

0.63±0.19

0.75±0.12

PMG-20/Silica

2.69±0.5

119±16

2.01±0.4

2.23±0.06

* Values taken are average of two experiments. Dynamic Mechanical Properties. The dynamic mechanical properties of the PMY, PMG-20, PMY/Silica and PMG-20/Silica compounds were studied to see the performance of the nanocomposites. The temperature dependence of storage modulus (E’) is presented in Figure S12 (SI) and the results are summarized in Table S5(SI). The increase in E’ of PMG-20/Silica compared to PMY/Silica nanocomposite at a given temperature was ascribed to larger interfacial interaction between silica and PMG-20 polymer. Tg and tan δ values of the PMY, PMG-20, PMG-20/Silica and PMY/Silica nanocomposites are summarized in Table S6(SI).The two neat elastomers exhibited 33 ACS Paragon Plus Environment


difference in the Tg and tan δ at Tg in the temperature dependence curve of tan δ. The curves changed notably after silica was added. From the temperature dependence of tan δ curve in Figure 9, it was clear that the Tg value increased significantly in the PMG20/Silica nanocomposite due to the presence of GMA in PMG-20 polymer in comparison to neat PMY polymer nanocomposite. The increase in Tg (-16 °C) of PMG-20/Silica nanocomposite suggests stronger rubber-filler interaction, indicating in reduction of the mobility of the polymer chains and improved dispersion of silica.48,49 The interfacial interaction in PMY/Silica compound is lower due to the absence of GMA part leading to a poorer silica dispersion and lowering of Tg (-46 °C). So, the DMA results manifest the improved reinforcement efficiency on the dispersion of silica and interfacial interaction between PMG-20 polymer and silica, showing the effect of epoxy groups in the performance of the nanocomposite.

Figure 9. Tan δ vs temperature plot of PMY, PMG-20, PMY/Silica and PMG-20/Silica nanocomposites. 34 ACS Paragon Plus Environment

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The tan δ values at -20 °C, -10 °C and 0 °C (indirect measure of winter traction, ice traction and wet skid resistance respectively) for the samples were evaluated (Table 6). The inset plot with enlarged temperature region is shown in tan δ curve (Figure 9). The PMG-20/Silica vulcanizate displayed good wet skid resistance (higher tan δ at 0 °C) as well as higher winter and ice traction (higher tan δ at -20 and -10 °C) compared to PMY/Silica vulcanizate. The high tan δ value at 0 °C indicates good wet skid resistance of tires.50-52 Considering both the wet grip and traction, PMG-20/Silica shows an attractive choice of this elastomer for tire application. Table 6. Dynamic Mechanical Properties of the Elastomer Vulcanizates. Sample Properties

PMY/Silica 0.31

Winter traction (tan δ at -20 °C)

PMG-20/Silica 0.98

Ice traction (tan δ at -10 °C)

0.14

0.86

Wet skid resistance (tan δ at 0 °C)

0.07

0.60

Swelling Behavior. In order to further investigate the effect of epoxy groups on the reinforcement of silica and the crosslinking between PMG elastomer and silica, the swelling ratio of the prepared unfilled and silica filled nanocomposites was measured. For the calculations of the swelling ratio Q (using Eq. (1)), the density of the solvent (ρs) and density of the elastomer (ρr) were used. The swelling ratio of unfilled PMY (Q=3.66) is higher than PMG-20 (Q=2.93). Due to additional crosslinking in PMG-20 elastomer because of the epoxy group the swelling value of the unfilled PMG-20 vulcanizate was lower than the unfilled PMY vulcanizate. The swelling ratio obtained for PMG-20/Silica compound (Q=2.50) was lower than that of PMY/Silica vulcanizate (Q=3.18). So, the 35 ACS Paragon Plus Environment


values effectively signify that the presence of epoxy group in PMG-20 elastomer provides effective crosslink points and stronger interfacial interaction with the surface of silica. As the GMA content increased in the polymer, the swelling ratio decreased and the crosslink density increased. Morphological Analysis. The morphology of the nanocomposite was observed using SEM analysis. Cryo fractured sample surface was taken to observe the dispersion of filler in the elastomer matrix. From the SEM images, the white phase represents the dispersed silica particles in the nanocomposite. Figure 10a-b displays the SEM images of PMY/Silica filled sample at low and high magnifications. Both the images show few agglomerations of silica particle and void space, indicating poor dispersion of silica in PMY/Silica compound. Figure 10c-d portrays the SEM images of PMG-20/Silica filled sample at low and high magnifications. The presence of GMA in PMG-20/Silica polymer nanocomposite remarkably leads to (Figure 10c and 10d) good uniform dispersion of silica preventing the aggregation of silica particles. The improved dispersion was due to interfacial interaction and efficient reinforcement of silica on PMG-20 elastomer matrix in the nanocomposite. The dispersion of silica in the nanocomposite will obviously vary with different composition of the copolymer. With greater percentage of GMA, the silica dispersion will be much better due to interaction with the epoxy group of GMA moiety. This is clear from Figure 10c-d. Figure S13 (SI) shows with higher fraction of GMA in the copolymer shows better dispersion in the resulting nanocomposite.


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Figure 10. SEM images of (a) and (b) PMY/Silica nanocomposite (at low and high magnification). (c) and (d) PMG-20/Silica nanocomposite (at low and high magnification). CONCLUSIONS.

The

present

work fabricates a

cross-linkable epoxy group-

functionalized biobased elastomer poly (myrcene-co-glycidyl methacrylate) (PMG) via redox emulsion polymerization and its silica nanocomposite for the first time. For comparison purpose, both PMY and PMY/Silica composite were discussed. The copolymer with 90 wt % β-myrcene and 10 wt % of GMA, poly (MY90GMA10) was considered for detailed structural characterization study. It displayed a highest molecular weight of 105,870 Da and a subzero glass transition temperature of −48 °C. The assignment of each peak by NMR supports the formation of the copolymer structure. A 37 ACS Paragon Plus Environment


set of 2-D NMR experiments helped to entirely assign the resonances of the spectra of coupled nuclear spins in the copolymers. The reactivity ratio shows azeotropic copolymerization behavior. Poly (MY80GMA20) copolymer and PMY homopolymer was compounded and vulcanized using conventional recipe and rubber processing methods. The presence of epoxy group in PMG/Silica vulcanizate effectively improved dispersion of silica in the matrix and performance because of the covalent bonding interaction between silica and PMG elastomer. These are shown by SEM, mechanical properties, swelling data and dynamic mechanical properties. FTIR result supported the improved interfacial interaction in PMG/Silica nanocomposite, due to ring-opening reaction between epoxy group of PMG elastomer and hydroxyl group of silica. PMG-20/Silica nanocomposite demonstrated better wet skid resistance, greater traction, and better mechanical properties than the PMY/Silica nanocomposite making it a good choice for some engineering applications. Therefore, the emulsion polymerized PMG polymer can be a solution to improve the performances of silica tire tread composite and have a great potential in the tire industry. Because of the incorporation of the polar group, these types of functionalized elastomers could be the possible alternative for different synthetic elastomers, to display fine interaction with various functional fillers used in the rubber industry. ASSOCIATED CONTENT

Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: 38 ACS Paragon Plus Environment

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Picture of synthesized copolymers; Relation between the particle size and distribution of copolymer; Calculation and parameters of reactivity ratios of monomers; Theoretical solubility parameter calculation and group contribution; 1H NMR and MY, GMA, PGMA and PMY; 1H NMR and

13

C NMR spectra of

13

C NMR spectra of the poly (MY80GMA20)

copolymer; COSY and NOESY (1H-1H) spectrum of PGMA, PMY and poly (MY90GMA10) copolymer; TGA thermogram and X-ray diffractograms of homopolymers and copolymers; FTIR comparison spectra of PMG/silica nanocomposites; E’ vs temperature plot of PMY and PMG-20 silica nanocomposites; E’, Tan δ and Tg values of PMY, PMG20, PMY/Silica and PMG-20/Silica nanocomposites; SEM images of PMG-10/Silica and PMG-20/Silica nanocomposite. AUTHOR INFORMATION Corresponding Author *Tel: +91-3222-283180. Fax: +91-3222-220312. E-mail: [email protected]; [email protected]. ORCID Anil K. Bhowmick: 0000-0002-8229-5353 Notes The authors declare no competing financial interest. Acknowledgements The authors would like to thank IIT Kharagpur for providing the necessary facilities. Pranabesh Sahu is thankful to CSIR (HRDG), New Delhi, for providing financial assistantship in the form of a junior research fellowship (Ref. No: 20/12/2015(ii) EU-V).



AKB acknowledges the partial support of Uchchatar Avishkar Yojana (UAY), MHRD, New Delhi and Central Research, Bridgestone Corporation, Japan.

REFERENCES 1) Zhu, Y.; Romain, C.; Williams, C. K. Sustainable polymers from renewable resources. Nature 2016, 540, 354−362. 2) Gandini, A. Polymers from Renewable Resources: A Challenge for the Future of Macromolecular Materials. Macromolecules 2008, 41, 9491−9504. 3) Coates, G. W.; Hillmyer, M. A. Polymers from Renewable Resources. Macromolecules 2009, 42, 7987−7989. 4) Gandini, A.; Lacerda, T. M.; Carvalho, A. J.; Trovatti, E. Progress of polymers from renewable resources: furans, vegetable oils, and polysaccharides. Chem. Rev. 2016, 116, 1637−1669. 5) Llevot, A.; Dannecker, P.; Czapiewski von, M.; Over, L. C.; Söyler, Z.; Meier, M. A. R. Renewability is not enough: Recent Advances in the Sustainable Synthesis of Biomass-Derived Monomers and Polymers. Chem. Eur. J. 2016, 22, 11510 11521. 6) Wang, Z.; Yuan, L.; Tang, C. Sustainable Elastomers from Renewable Biomass. Acc. Chem. Res., 2017, 50, 1762–1773. 7) Bhowmick, A. K.; Stephens, H. L.; Handbook of Elastomers, Marcel Dekker AG, New York, 2001, ch. 2.


Page 40 of 48

Page 41 of 48 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


8) Firdaus, M.; Montero de Espinosa, L.; Meier, M. A. R. Terpene-Based Renewable Monomers and Polymers via Thiol-Ene Additions. Macromolecules 2011, 44, 7253−7262. 9) Wilbon, P. A.; Chu, F.; Tang, C. Progress in Renewable Polymers from Natural Terpenes, Terpenoids and Rosin, Macromol. Rapid Commun. 2013, 34, 8-37. 10) Bolton, J. M.; Hillmyer, M. A.; Hoye, T. R. Sustainable Thermoplastic Elastomers from Terpene-Derived Monomers. ACS Macro Lett. 2014, 3, 717−720. 11) Winnacker, M.; Rieger, B. Recent Progress in Sustainable Polymers Obtained from

Cyclic

Terpenes:

Synthesis,

Properties,

and

Application

Potential.

ChemSusChem 2015, 8, 2455−2471. 12) Breitmaier, E. Terpenes: Flavors, Fragrances, Pharmaca, Pheromones; WileyVCH Verlag GmbH: Weinheim, 2006. 13) Biermann, U.; Bornscheuer, U.; Meier, M. A. R.; Metzger, J. O.; Schäfer, H. J. Oils and Fats as Renewable Raw Materials in Chemistry. Angew. Chem. Int. Ed. 2011, 50, 3854−3871. 14) Surburg, H.; Panten, J. Common Fragrance and Flavor Materials, 5th ed., WileyVCH, Weinheim, 2006. 15) Behr, A.; Johnen, L. Myrcene as a Natural Base Chemical in Sustainable Chemistry ChemSusChem 2009, 2, 1072 – 1095. 16) Sarkar, P.; Bhowmick, A. K. Synthesis, characterization and properties of a biobased elastomer: polymyrcene. RSC Adv. 2014, 4, 61343−61354.



Page 42 of 48

17) Sarkar, P.; Bhowmick, A. K. Green Approach toward Sustainable Polymer: Synthesis and Characterization of Poly-(myrcene-co-dibutyl itaconate). ACS Sustainable Chem. Eng. 2016, 4, 2129−2141. 18) Sarkar, P.; Bhowmick, A. K. Terpene Based Sustainable Elastomer for Low Rolling

Resistance

and

Improved

Wet

Grip

Application:

Synthesis,

Characterization and Properties of Poly (styrene-co-myrcene). ACS Sustainable Chem. Eng. 2016, 4, 5462−5474. 19) Sarkar, P.; Bhowmick, A. K. Terpene Based Sustainable Methacrylate Copolymer Series

by

Emulsion

Polymerization:

Synthesis

and

Structure-Property

Relationship. J. Polym. Sci., Part A: Polym. Chem. 2017, 55, 2639–2649. 20) Brar, A. S.; Dutta, K. Acrylonitrile and Glycidyl Methacrylate Copolymers: Nuclear Magnetic Resonance Characterization. Macromolecules 1998, 31, 4695–4702. 21) Brar, A. S.; Pradhan. D. R. Investigation of microstructure of the acrylonitrilestyrene-glycidyl methacrylate terpolymers by 1D and 2D NMR spectroscopy. Journal of Applied Polymer Science 2003, 89, 1779–1790. 22) Brar, A. S.; Yadav, A. Microstructure of Glycidylmethacrylate/Vinyl Acetate Copolymers by Two-Dimensional Nuclear Magnetic Resonance Spectroscopy. J. Polym. Sci., Part A: Polym. Chem. 2001, 39, 4051–4060. 23) Dhal, P. K.; Babu, G. N. Glycidyl Methacrylate-tert-Butyl Acrylate Copolymers: Synthesis, Characterization, and Thermal Studies. J. Polym. Sci., Part A: Polym. Chem. 1984, 22, 1817-1829.


Page 43 of 48 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


24) Höhne, S.; Uhlmann, P. Synthesis of Functional Block Copolymers and Terpolymers Containing Polyglycidyl Methacrylate Blocks. J. Polym. Sci., Part A: Polym. Chem. 2014, 53, 675–684. 25) El-Aasser, M.; Lovell, P. Emulsion Polymerization and Emulsion Polymers; John Wiley & Sons, Inc.: Chichester, 1997. 26) Gilbert, R. G. Emulsion Polymerization: A Mechanistic Approach; Academic Press: London, 1995. 27) Lei, W.; Zhou, X.; Russell, T. P.; Hua, K.; Yang, X.; Qiao, H.; Wang, W.; Li, F.; Wang, R.; Zhang, L. High performance bio-based elastomers: energy efficient, sustainable materials for tires. J. Mater. Chem. A, 2016, 4, 13058-13062. 28) Edwards, D. C. Polymer-filler interactions in rubber reinforcement. J. Mater. Sci. 1990, 25, 4175–4185. 29) Wang, M. J. Effect of polymer-filler and filler-filler interactions on dynamic properties of filled vulcanizates. Rubber Chem. Technol. 1998, 71, 520-589. 30) Hamed, G. R. Reinforcement of Rubber. Rubber Chem. Technol. 2000, 73, 524533. 31) Qiao, H.; Xu, W.; Chao, M.; Liu, J.; Lei, W.; Zhou, X.; Wang, R.; Zhang, L. Preparation and Performance of Silica/Epoxy Group-Functionalized Biobased Elastomer Nanocomposite. Ind. Eng. Chem. Res. 2017, 56, 881−889. 32) Qiao, H.; Wang, R.; Yao, H.; Wu, X.; Lei, W.; Zhou, X.; Hu, X.; Zhang. L. Design and preparation of natural layered silicate/bio-based elastomer nanocomposites with improved dispersion and interfacial interaction. Polymer 2015, 79, 1-11.



Page 44 of 48

33) Flory, P. J.; Rehner, J., Jr. Statistical Mechanics of Cross-Linked Polymer Networks II. Swelling. J. Chem. Phys. 1943, 11, 521−526. 34) Chern, C. S. Emulsion Polymerization Mechanisms and Kinetics. Prog. Polym. Sci. 2006, 31, 443−486. 35) Thickett, S. C.; Gilbert, R. G. Emulsion Polymerization: State of the Art in Kinetics and Mechanisms. Polymer 2007, 48, 6965−6991. 36) Krevelen, D. W. V.; Nijenhuis, K. T. Properties of Polymers: Their Correlation with Chemical Structure. Their Numerical Estimation and Prediction from Additive Group Contributions, 4th ed.; Elsevier: Amsterdam, 2009; Chapter 7. 37) Fineman, M.; Ross, S. D. Linear Method for Determining Monomer Reactivity Ratios in Copolymerization. J. Polym. Sci. 1950, 5, 259−262. 38) Kelen, T.; Tüdös, F. Analysis of the Linear Methods for Determining Copolymerization Reactivity Ratios. I. A New Improved Linear Graphic Method. J. Macromol. Sci., Chem. 1975, 9, 1−27. 39) Odian, G. Principles of Polymerizations, 4th ed.; Wiley Interscience: USA, 2004; Chapter 6. 40) Lovell, P. A.; El-Aasser, M. S. Emulsion Polymerization and Emulsion Polymers edited by John Wiley and Sons Ltd. Chapter 1. 41) Potts, J. R.; Shankar, O.; Du, L.; Ruoff, R. S. Processing−Morphology−Property Relationships

and

Composite

Theory

Analysis

of

Reduced

Graphene

Oxide/Natural Rubber Nanocomposites. Macromolecules 2012, 45, 6045−6055. 42) Edwards D. Polymer–filler interactions in rubber reinforcement. J. Mater. Sci. 1990, 25, 4175–4185.


Page 45 of 48 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


43) Zheng, H.; Zhang, Y.; Peng, Z.; Zhang, Y. Influence of clay modification on the structure and mechanical properties of EPDM/montmorillonite nanocomposites. Polym. Test. 2004, 23, 217-223. 44) Jia, Q.X.; Wu, Y.P.; Wang, Y.Q.; Lu, M.; Yang, J.; Zhang, L.Q. Organic interfacial tailoring of styrene butadiene rubber–clay nanocomposites prepared by latex compounding method. J. Appl. Polym. Sci. 2007, 103, 1826-1833. 45) Gui, Y.; Zheng, J.; Ye, X.; Han, D.; Xi, M.; Zhang, L. Preparation and performance of silica/SBR master batches with high silica loading by latex compounding method. Composites, Part B, 2016, 85, 130−139. 46) Qu, L.; Yu, G.; Xie, X.; Wang, L.; Li, J.; Zhao, Q. Effect of silane coupling agent on filler and rubber interaction of silica reinforced solution styrene butadiene rubber. Polym. Compos. 2013, 34, 1575−1582. 47) Yan, H.; Tian, G.; Sun, K.; Zhang, Y.; Zhang, Y. Effect of silane coupling agent on the polymer-filler interaction and mechanical properties of silica-filled NR. J. Polym. Sci., Part B: Polym. Phys. 2005, 43, 573−584. 48) Wang, M. The role of filler networking in dynamic properties of filled rubber. Rubber Chem.Technol. 1999, 72, 430–448. 49) Wang, M. Effect of polymer–filler and filler–filler interactions on dynamic properties of filled vulcanizates. Rubber Chem. Technol. 1998, 71, 520–589. 50) Wang, M.; Kutsovsky, Y. Effect of fillers on wet skid resistance of tires. Part I: water lubrication vs. filler-elastomer interactions. Rubber Chem. Technol. 2008, 81, 552–575.



51) Takino, H.; Nakayama, R.; Yamada, Y.; Kohjiya, S.; Matsuo, T. Viscoelastic Properties of Elastomers and Tire Wet Skid Resistance. Rubber Chem. Technol. 1997, 70, 584-594. 52) Wang, Q.; Liu, J.; Cui, Q.; Xiao, X. Effect of Elastomer Nanoparticles on Improving the Wet Skid Resistance of SBR/NR Composites. Rubber Chem. Technol. 2016, 89, 262-271.


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For Table of Contents use only

Synopsis: A multi-functional bio-based elastomer having excellent mechanical properties was developed via a sustainable method for improved silica dispersion and reinforcement



For Table of Contents use only

Synopsis: A multi-functional bio-based elastomer having excellent mechanical properties was developed via a sustainable method for improved silica dispersion and reinforcement

ACS Paragon Plus Environment

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Design of a Molecular Architecture via a Green Route for an Improved

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