NR Ternary Thermoplastic Vulcanizates

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Bio-Based PLA/NR-PMMA/NR Ternary TPVs with Balanced Stiffness and Toughness: “Soft-Hard” Core-Shell Continuous Rubber Phase, In-Situ Compatibilization and Properties Yukun Chen, Wentao Wang, Daosheng Yuan, Chuanhui Xu, Liming Cao, and Xingquan Liang ACS Sustainable Chem. Eng., Just Accepted Manuscript • DOI: 10.1021/ acssuschemeng.8b00267 • Publication Date (Web): 27 Mar 2018 Downloaded from http://pubs.acs.org on March 27, 2018

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ACS Sustainable Chemistry & Engineering

Bio-Based PLA/NR-PMMA/NR Ternary TPVs with Balanced Stiffness and Toughness: “Soft-Hard” Core-Shell Continuous Rubber Phase, In-Situ Compatibilization and Properties Yukun Chen1,*, WentaoWang1, Daosheng Yuan1, Chuanhui Xu 1,2,*, Liming Cao1, Xingquan Liang2 1

Lab of Advanced Elastomer, South China University of Technology, Wushan RD., Tianhe

District, Guangzhou, 510640, China 2

School of Chemistry and Chemical Engineering, Guangxi University, No. 100, Daxuedong Road,

Xixiangtang District, Nanning, 530004, China Corresponding Author: Yukun Chen [email protected] and Chuanhui Xu [email protected] ABSTRACT: Stiffness and toughness are two mutually exclusive attributes of polymer materials that contribute to significant improvement in impact strength is usually accompanied with a reduction in tensile strength. In this study, ternary thermoplastic vulcanizates (TPVs) consisting of poly (lactic acid) (PLA), poly (methyl methacrylate) grafted natural rubber (NR-PMMA) and natural rubber (NR) with balanced stiffness and toughness were successfully prepared via peroxide-induced dynamic vulcanization. With 10 wt % of NR and NR-PMMA, the PLA/NR-PMMA/NR ternary TPV displayed an enhanced yield stress of 41.7MPa (only 38% loss compared to neat PLA) and a significantly higher impact strength of 91.30 kJ/m2 (nearly 32 times that of neat PLA). The in-situ compatibilization between PLA and rubber phases was confirmed by Fourier transform infrared spectroscopy (FT-IR). Interfacial, rheological and calorimetric measurements confirmed that the NR was encapsulated by NR-PMMA in the PLA phase. It was found that the flexibility of the “soft” NR core and outer “hard” NR-PMMA shell with excellent PLA/rubber interfacial adhesion are responsible for the super toughness and considerable tensile strength of the PLA/NR-PMMA/NR ternary TPVs. Keywords: dynamic vulcanization, super-tough, balanced stiffness and toughness, “soft-hard” core-shell structure INTRODUCTION Given the growing concerns on sustainability and environment, bio-based polymer materials derived from renewable resources are increasingly utilized as alternatives to traditional petroleum-based materials.1,2 Polylactide (PLA), with excellent biocompatibility, bioresorbability , biodegradability as well as outstanding mechanical properties, is definitely the most promising bio-based polymer at present.3,4 However, the brittleness of PLA, to a great extent, limits its applications.5 Until now, many toughening strategies for PLA such as copolymerization,6 plasticization,7,8 blending9,10 and nanocomposition11 have been extensively studied. Blending with elastomeric component is an efficient way to achieve the appreciable toughness for PLA, especially when some physical or chemical interactions are generated at the interface between rubber and PLA.12,13 Recently, considerable attentions have been given to dynamic vulcanization of PLA and rubbers, where the in-situ compatibilization reactions can be occurred. Wang et al.14 reported a super toughened PLA/unsaturated aliphatic polyester elastomer (UPE) thermoplastic vulcanizate (TPV) prepared via peroxide-induced dynamic vulcanization. 1

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The TPV manifested a quasi co-continuous morphology with vulcanized UPE compactly dispersed in PLA matrix. Ma et al.15 prepared a TPV with excellent properties by dynamic cross-linking of PLA and ethylene-co-vinyl acetate (EVA). They found that the morphology transition from typical “sea−island” morphology to dual-continuous network-like structure occurred in the TPV with increased free radical initiators. In addition to this, thylene/n-butylacrylate/glycidyl methacrylate (EBA-GMA),16,17 poly(ethylene−glycidyl 18 19 methacrylate) (EGMA), crosslinked polyurethane (CPU), thermoplastic elastomer 20 poly(ester−amide) (PEA), and bio-based polyester elastomer (BPE) 21 were also reported to toughen PLA. Although the above-mentioned elastomers worked well in toughening PLA, most of them are synthetic petroleum-based materials. In contrast, natural rubber (NR), known as a “green” and renewable material from rubber trees, is a suitable substitute to toughen PLA without impairing its bio-based attribute.22,23 Bitiniset al.24,25 have prepared the PLA/NR physical blend, which got the best comprehensive performances at 10 wt % of NR. However, the improvement of toughness was still not satisfied due to the great polarity difference between PLA and NR. Hence, modified natural rubbers (NR), such as NR-g-(glycidyl methacrylate) (GMA),26 NR-poly(butyl acrylate) (PBA),27 PLA-g-NR,28 PLA-NR-PLA29 and so on, were utilized as a toughener or compatibilizer to strengthen the PLA/NR system through physical blending. Recently, Wu et al.30 have reported a dynamically vulcanized PLA/NR-GMA TPV systems with a significantly improved notched impact strength of 73.4kJ/m2, approximately a 26- increased. Unfortunately the tensile strength suffered a nearly 53% loss. This is a quite unpleasant empirical law that the incorporation of rubbers with much lower modulus usually result in an inevitable reduction in tensile strength for polymer blends. Therefore, in this sense, to achieve a balance between stiffness and toughness of PLA/rubber system is always pragmatic and meaningful for PLA-toughening industry. In our previous work, we also demonstrated a super toughened bio-based PLA/NR TPV with novel continuous crosslinked rubber phase.31 We obtained a significant high impact strength which was up to 58.3kJ/m2 with only 35% rubber content. However, the tensile strength was declined sharply from 60 MPa to 22.5 MPa, about 62.5% loss.31,32 To tailor the balance between the stiffness and toughness of above PLA/NR TPVs, in this work, we incorporated NR and poly (methyl methacrylate) grafted natural rubber (NR-PMMA)33 into PLA and prepared dynamically vulcanized PLA/NR-PMMA/NR ternary TPVs. Because of the grafted PMMA, the modulus of NR-PMMA was much higher than that of NR. Considering the polarity differences among PLA, NR and NR-PMMA, the “soft” NR was inclined to be encapsulated by the “hard” NR-PMMA, forming a “soft-hard” core-shell continuous rubber structure in PLA phase. Due to this specific structure, we achieved a marvelous balance in stiffness/toughness performance of PLA/NR-PMMA/NR TPVs. The mechanical properties, interfacial compatibility and crystallization behavior were thoroughly investigated and discussed based on this phase structure, and also the toughening mechanism was established according to the relationship between the structure and property. EXPERIMENTAL SECTION Materials and sample preparation Polylactide (PLA), REVODE101 (Zhejiang HisunBiomaterials Co., Ltd), MI (190°C , 2.16 kg) =5~8 g/10 min, weight average molecular weight (Mw)=150000 g/mol, specific gravity=1.25 g/cm3; nature rubber (NR), SMRCV60 (Guangzhou Rubber Industry Research Institute, China), 2


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Mooney viscosity ML(1+4)100°C =60 ± 5, specific gravity=0.93 g/cm3; Thirty-five weight percent poly (methyl methacrylate) grafted natural rubber, MG35 (Guangzhou Rubber Industry Research Institute, China), Mooney viscosity ML(1+4)100°C=100 ± 5, specific gravity=0.98 g/cm3; Dicumyl peroxide (DCP), Sinopharm Chemical Reagent Co. Ltd (China), purified by anhydrous alcohol recrystallization before use. Irganox 1010, antioxidant, industrial grade and obtained on the open market, which is used to prevent the degradation of PLA during processing. The other chemicals were used as received. The PLA/NR-MMA/NR TPVs were prepared in a Haake Rheocord 90 at 150°C and a rotor speed of 60 rpm. Before used, PLA was dried for 6 h in a vacuum oven at 80°C. The dried PLA with quantitative Irganox 1010 was firstly shear-melted for 3 mins and then the masticated NR and NR-PMMA were added. After reaching the stable torque, DCP was added to induce the dynamic vulcanization. Four minutes later, the TPV was removed from the mixer and cooled down to the room temperature. Next, the large blocks of TPV were chopped into small granules for injection molding in a TTI-160F (Welltec Machinery &Equipment Co. Ltd., China). The temperature profile of the injection barrels was 165/170/170/175 °C from the first heating zone to nozzle, respectively. The injection pressure was 35 MPa. For all of the TPVs, the weight ratio of DCP was maintained at 2.0 wt% of the total rubber content. The concentration of Irganox 1010 was fixed at 0.2 wt% of (PLA + NR + NR-PMMA). For convenience, the sample codes were defined according to the PLA/NR-PMMA/NR ratio, as shown in Table 1. Table 1 Sample compositions and the corresponding codes. PLA/NR-PMMA/NR

Composition (wt.%)

Sample code

TPV

80/0/20

D80/0/20

TPV

80/5/15

D80/5/15

TPV

80/10/10

D80/10/10

TPV

80/15/5

D80/15/5

TPV

80/20/0

D80/20/0

blend

80/10/10

B80/10/10

Mechanical properties The tensile properties of the dumbbell shaped specimens were measured in accordance with ISO 527 at room temperature using a universal testing instrument (Shimadzu AG-1, 10kN, Japan), tensile speed was 50 mm/min. The notched Izod impact tests were carried out according to ISO 180 using an impact test machine (ZWICK5331, German, Zwick/Roell) at room temperature. The average value was calculated from at least 5 test specimens. Instrumented notched impact tests were conducted on a FRACTOVIS PLUS manufactured by CEAST. Tests are performed with standard specimens measuring 63mm×10mm×4mm equipped with 45 V-shaped notches of 2 mm depth and 0.25 mm root radius at the tip of the notch, the rate of fracture 3.5 m/s. Contact angle measurement Contact angles were measured by the sessile drop technique using an apparatus model OCA 15 PLUS, DATAPHYSICS. Contact angles measurement of a given material was carried out at 3



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least 10 times at different positions. The final contact angle was the average value of 10 results. Fourier transform infrared spectroscopy (FT-IR) The FT-IR absorption spectra were recorded by a Tensor 27 Spectrometer (Bruker, Germany) in the wave number range from 400 to 4000 cm−1 with a resolution of 4 cm-1 and 32 scans. Neat PLA, NR-PMMA and NR were directly tested under attenuated total reflectance (ATR) model. For the PLA/NR-PMMA/NR TPVs, their samples were first extracted with dichloromethane (DCM) at ambient temperature for 72 h to selectively remove the free PLA thoroughly, and then oven-dried in a vacuum at 60 °C to eliminate the residual solvent and moisture. Finally, the residue rubber components were compressed into disks for the FT-IR test under ATR model. Crosslink density of the TPVs The weighed pieces of TPVs were immersed in toluene at room temperature for 5d until reaching swelling equilibrium. Then, the swollen samples were blotted with tissue paper to remove the solvent and weighed again. Vr, the volume fraction of rubber swollen in the sample gel, was used to represent the apparent cross-link density, which was calculated by the following equation:34-36 1 Vr = ρ m2 1+ ( − 1) × r m1 αρ s (1) where m1 and m2 are the mass of the sample before and after swollen; ρ r and

ρs are the

rubber (NR: 0.93 g/cm3, NR-PMMA: 0.98 g/cm3) and toluene density (ρs = 0.865 g/cm3), respectively; α is the mass fraction of rubber phase in the sample. Analysis of crystallization Differential scanning calorimetry measurements were performed in a NETZSCH DSC 204 under a nitrogen atmosphere. For each test, samples of ~5 mg sealed in aluminum pans were heated from room temperature to 190°C at a rate of 30°C /min and then held at 190°C for 5 min to eliminate the thermal history. Afterwards, the samples were cooled to 20°C at a rate of 20°C /min, holding about 3 min at 20°C, and reheated to 190°C at a heating rate of 20°C/min. The crystallinity of PLA ( X c ) in the samples was estimated from the second heating cycle using the following equation:14,15

Xc =

∆H m − ∆H c × 100% ϕ∆H m0

(2)

where ∆H m and ∆Hc are the enthalpies of melting and cold crystallization during heating, respectively; ∆H m0 is the enthalpy assuming 100% crystalline PLA homopolymer (93.7 J/g), and

ϕ is the weight fraction of PLA component in the sample. Scanning electron microscopy (SEM) The morphology of the TPVs was observed using Merlin field emission scanning electron microscopy (FE-SEM, Carl Zeiss). The morphologies of fracture surfaces were obtained from notched Izod impact test and the morphologies of fracture surface of tensile samples were obtained from tensile test. In order to observe the crosslinked NR phase, blends were subjected to 4


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dichloromethane wash (ambient temperature) to remove the PLA phase at the surface and then dried sufficiently. All the surfaces of samples were sputter coated with a thin layer of gold prior to morphological observation. RESULTS AND DISCUSSION Mechanical properties of TPVs The impact strengths of neat PLA, PLA/NR-PMMA/NR ternary TPVs and corresponding PLA/NR and PLA/NR-PMMA binary TPVs are shown in Figure 1a. In corporation of 20 wt% NR, the impact strength of PLA was only increased from 2.9 kJ/m2 to 7.1 kJ/m2 (D80/0/20). When NR was replaced by NR-PMMA, the impact strength of D80/20/0 was increased to 31.9 kJ/m2, nearly 11 times that of neat PLA. Interestingly, when NR and NR-PMMA were both used to toughen PLA, a significant synergistic effect was achieved by changing the ratios of NR/NR-PMMA. As clearly seen, PLA/NR-PMMA/NR ternary TPVs showed a remarkable toughening improvement that D80/10/10 obtained the highest impact strength of 91.3 kJ/m2, nearly 32 times that of neat PLA. However, B80/10/10, a simple blend without dynamic vulcanization, showed poor impact strength of only 14.4 kJ/m2. This revealed that the dynamic vulcanization had an important effect on the mechanical properties of the TPV. It was worth noticing that all of the test specimens of PLA/NR-PMMA/NR ternary TPVs (D80/5/15, D80/10/10 and D80/15/5) were not completely fractured during the impact process, as shown in the inset of Figure 1a. The actual values obtained from the tests were represented in sparse column, and the final impact strength values were the ones added with the contributions from the unbroken parts, which were much higher than that of D80/0/20 and D80/20/0.

Figure 1. Mechanical properties of neat PLA and PLA/NR-PMMA/NR TPVs: (a) Izod impact strength; (b) stress-strain curves; (c) yield stress and (d) Young's modulus 5



Figure 1b shows the tensile behaviors of neat PLA, D80/0/20, D80/20/0, D80/10/10 and B10/10/10. The stress−strain curve of neat PLA showed a typical brittle fracture with a tensile strength of 67.5 MPa (elongation 7.5%). Because of the soft nature of NR, the stress–strain curve of D80/0/20 exhibited the lowest yield stress of 36.4 MPa (almost half loss). As for D80/20/0, the better compatibility of NR-PMMA/PLA interface saved considerable yield stress of PLA phase. It exhibited characteristic nature of ductile fracture with apparent yielding in the stress−strain curve and showed a higher yield stress of 48.5 MPa. As expected, D80/10/10 had a typical tensile behavior of toughened polymer blends with a moderate yield stress of 41.7 MPa, while B80/10/10 showed an inferior tensile behavior with a yield stress of 38.1 MPa. We also summarized the yield stress and the corresponding Young's modulus of the samples in Figure 1c and d. Because the modulus and the compatibility of NR-PMMA were higher than those of NR, the yield stress and Young's modulus of the TPVs were improved with the increasing NR-PMMA content. Therefore, the above experimental results demonstrated an exciting fact that the toughness of PLA could be significantly improved without much decline in its strength by tailoring NR-PMMA/NR ratios. By this way, a bio-based PLA TPV with good balance between stiffness and toughness can be designed. Co-continuous phase structure of TPVs

Figure 2. SEM micrographs of cryo-fracture surface etched by dichloromethane: (a) D80/0/20; (b) D80/5/15; (c) D80/10/10; (d) D80/15/5; (e) D80/20/0 In our previous study of PLA/NR system, we found that the cross-linked NR phase in PLA/NR TPV was a specific continuous structure rather than dispersed particles.31,32 Interestingly, the similar phase morphology structure was also found in this PLA/NR-PMMA/NR ternary system. As shown in Figure 2, after the PLA phase in the surface layer was removed, the residual rubber phase was a continuous structure. Because of the continuous phase of PLA, TPVs formed a co-continuous phase structure, which was similar to the “Steel Reinforced Concrete Structure”.31,32,34 This was attributed to that melt strengths of NR and NR-PMMA were much higher than that of PLA.31,33 As a result, the rubber phase was unable to be broken up effectively, and thus remained an IPN structure during dynamic vulcanization. According to our previous study, the rubber phase was partial continuous rather than entire continuous in PLA phase, which was not influence the molding of TPV. More evidences for the continuous rather phase in PLA can be found in the swelling experiments, as shown in Figure S1. B80/10/10 also exhibited similar rubber phase morphology, which was shown in Figure S2. 6


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Discussion on the location of NR-PMMA and NR We used interfacial tensions of polymer pairs in PLA/NR-PMMA/NR ternary system to discuss the possible location of NR-PMMA and NR, which can be determined through contact angles experiments (see Figure S3 and Table S1). As shown in Table 1, the interfacial tension of PLA/NR pair was 3.75 mN/m, which was higher than the 2.22 mN/m of PLA/NR-PMMA pair. The most important fact was that the interfacial tension of PLA/NR-PMMA pair was the lowest, 0.29 mN/m, which strongly suggested that the interaction between PLA and NR-PMMA was much stronger than that between PLA and NR. Therefore, when the above three polymers were blended together, what most likely to happen was that NR-PMMA would locate at the interfacial region between the PLA phase and the NR phase to minimize the phase interfacial free energies,37 forming a “soft-hard” core-shell rubber phase structure in the PLA continuous phase. Table 1 Interfacial tensions of possible polymer pairs Polymer pairs

Interfacial tension (mN/m)

PLA/NR-MMA

0.29

PLA/NR NR-MMA/NR

3.75 2.22

In-situ compatibilization during dynamic vulcanization

Figure 3. FT-IR absorption spectra of (a) the individual polymers and (b) the residues of dichloromethane extracted PLA/NR-PMMA/NR TPVs The in-situ reactive compatibilization at the plastic/rubber interface occurred during DCP-induced dynamic vulcanization was quantitatively investigated by FT-IR. As shown in Figure 3a, the characteristic absorption peak at 1750 cm-1 was corresponded to the stretching vibration of carbonyl groups (C=O) of PLA,17,21 and the absorption peak at 1375 cm-1 was assigned to the symmetric stretching deformation of methyl (−CH3) of NR.32,34 In the FT-IR spectra of NR-PMMA, it showed characteristic absorption peaks at 1375 and 1730 cm-1, which were attributed to the symmetric stretching deformation of methyl (−CH3) and the stretching vibration of carbonyl groups (C=O) of the PMMA segment,33 respectively. As shown in Figure 3b, the spectra of the DCM-extracted residues were almost parallel to that of pure NR and NR-PMMA,31,32 suggesting a complete elimination of the free PLA after DCM-extraction. As expected, the absorption peak represented the carbonyl vibration (C=O) of PLA was shifted from 1750 cm-1 to a higher wavenumber of 1760 cm-1, as visible for all of the residues of DCM-extracted TPVs. Moreover, a gradually increased absorption peak at about 1730 cm-1 which represented the carbonyl vibration (C=O) of PMMA segment was accompanied with the 7



increasing NR-PMMA content. This strongly suggested that PLA molecules were reacted with the rubber chains during dynamic vulcanization.

Figure 4. (a) Ratio of absorption peak area (A1760/A1375) of TPVs and (b) Vr of the rubber phase in TPVs The extent of the in-situ compatibilization at the plastic/rubber interface can be approximately described by the amount of grafted PLA per unit of rubber residues, as measured by the ratio of absorption peak area at 1750 cm-1 to that at 1375 cm-1 in Figure 4a.31,32 Although the PMMA segment grafted on the molecular chain of NR-PMMA led to a decreased number of the active points, the PLA molecular free radicals could still be reacted with the C=C double bonds of NR-PMMA radicals. Therefore, the chance of direct contacts between PLA and rubber phase was promoted with the increased NR-PMMA contents, which resulted in a strong and stable interface between PLA and rubber phases. We also showed the Vr of rubber phase in the TPVs in Figure 4b. The Vr of rubber phase was slightly increased with more NR-PMMA, which might be attributed to that the “hard shell” of NR-PMMA restricted the free swelling in the “soft core” of NR due to which the total swollen of rubber phase was restrained by the NR-PMMA shell, resulting in an apparently increased Vr.

Figure 5. Schemes of the possible in-situ compatibilization at interface between PLA and rubber phases and the final core-shell rubber phase structure 8


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Base on the above analysis, the in-situ reactive compatibilization between crosslinked rubber and PLA during dynamic vulcanization was proposed as schematically illustrated in Figure 5. During the melt blending process, the requirement of minimizing the surface free energy drove NR-PMMA to locate between PLA and NR phases. At the same time, the carboxyl groups of NR-PMMA induced a dipole force with the carbonyl group of PLA to further enhance the interactions between PLA and PMMA grafts of NR-PMMA.38 Decomposition of DCP generated massive free radicals, which not only induced the crosslinking of rubber phase but also initiated the in-situ compatibilization at the plastic/rubber interface. When PLA molecular free radicals met rubber radicals, graft copolymers at the plastic/rubber interface would be formed.18,19,30 As a result, NR was encapsulated by NR-PMMA, forming a “soft-hard” core-shell local continuous phase structure in the PLA phase.37 During the tensile tests, the TPVs exhibited ductile traits with stress whitening and necking through cold drawing.39 In order to further understand the interfacial adhesion in the TPVs with different NR-PMMA/NR ratios, the necking regions of the stretched samples were cryogenically fractured longitudinally and the SEM micrographs are shown in Figure 6. Numerous voids and rubber microfibrils resulted from debonding or internal cavitations can be observed in Figure 6b. As seen, the region was smooth in D80/0/20, showing a phase-separated morphology due to the poor interfacial adhesion between the PLA phase and the rubber microfibrils.20 As shown in Figure 6c, the surface of D80/10/10 was rougher compared to D80/0/20 and the voids and elastic microfibrils resulting from “pulling out” disappeared due to the improved interfacial adhesion. However, a few of rubber microfibrils could be observed in B80/10/10 in Figure S4, which suggested that the dynamic vulcanization played an important role in improving the compatibility between PLA and NR-PMMA. As shown in Figure 6d, the interface between the PLA phase and the rubber phase became indistinct and there were no rubber microfibrils oriented in the direction of stretching, which suggested a further enhanced plastic/rubber interfacial compatibility. The disappearance of the rubber microfibrils was attributed to a higher modulus of the crosslinked NR-PMMA. As a result, the excellent interfacial interactions between the “hard” NR-PMMA/NR shell and PLA phase contributed to the considerable strength of the ternary TPVs.

9



Figure 6. SEM micrographs of stretched PLA/NR-PMMA/NRTPVs obtained from a tensile bar after necking as schematically indicated in (a): (b) D80/0/20; (c) D80/10/10 and (d) D80/20/0 Crystallization behavior

Figure 7. (a) DSC melting curves and (b) crystallinity of neat PLA and PLA/NR-PMMA/NR TPVs The crystallization behaviors of neat PLA and PLA/NR-PMMA/NR TPVs were analyzed by DSC. As seen from Figure 7a, the glass transition temperature (Tg) of PLA phase in TPVs was slightly increased from 63.1°C to 63.5°C as the NR-PMMA/NR ratio increased from 0/20 (D80/0/20) to 10/10 (D80/10/10). When the NR-PMMA/NR ratio >10/10 (w/w), the Tg of PLA phase was not changed. This was attributed to that the interactions between NR-PMMA and PLA restricted the mobility of the PLA chains. Similar tendency can be found in the degree of crystallinity (Xc) in Figure 7b.The higher the NR-PMMA content in TPVs, the lower the Xc was. The Xc decreased little when the content of NR-PMMA was beyond 10wt%. Both the results of crystallization behavior and crystallinity illustrated that the plastic/rubber interface interaction in D80/10/10, D80/15/5 and D80/20/0 was similar. Furthermore, the relative rigid molecule chains of the NR-PMMA also hindered the regular arrangement of PLA molecular chains during the crystallization process. This resulted in an increased crystallization temperature (Tc) for all TPVs. 10


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Toughening mechanism

Figure 8. Instrumented impact measurements of neat PLA and PLA/NR-PMMA/NR TPVs: (a) load-displacement curves; (b) fracture energy and toughness index; (c) schematic picture to show impact process and the fracture surface As mentioned previously, the good balance between super toughness and considerable stiffness was the highlight of this PLA/NR-PMMA/NR ternary TPV system. The ternary TPVs showed higher energy consumption during the instrumental impact process when compared to the binary TPVs. Here, we exhibited the load-displacement curves of pure PLA and PLA/NR-PMMA/NR TPVs obtained in the instrumented notched impact test in Figure 8a. As seen, the neat PLA and D80/0/20 showed similar curves that the impact forces increased up to a maximum value and then quickly dropped to zero as a brittle crack propagated, implying that most energy was consumed in the initiation stage. Although the displacement of D80/0/20 in the initiation stage was slightly larger than that of neat PLA, the toughening effect was insufficient when 20 wt% NR added. When both NR and NR-PMMA was added, D80/5/15 showed a completely different fracture behavior. The impact force dropped slowly after the maximum value, implying that most energy was consumed in the crack-propagation stage while little energy was dissipated in the crack initiation stage. Moreover, the crack-propagation distinctly became higher and stable in D80/10/10, then decreased slightly in D80/15/5. This suggested that the highest toughness of D80/10/10 was due to the highest energy absorption during the entire fracture process. The crack-initiation energy, crack-propagation energy and toughness index of neat PLA and PLA/NR-PMMA/NR TPVs are shown in Figure 8b. The schematic of impact process and the fracture surfaces are illustrated in Figure 8c. Both crack initiation energy and crack propagation energy were very low in PLA and D80/0/20, showing typical brittle characteristics with toughness index of 0.19 and 0.30, respectively. Similarly, the crack initiation energy and crack propagation energy of the PLA/NR-PMMA binary TPV system were increased, and the former was slightly 11



larger than the latter. In addition, toughness index of D80/20/0 was 0.96, which was 5.1 times than that of neat PLA. On the contrary, the crack-initiation energy and crack-propagation energy of PLA/NR-PMMA/NR ternary TPVs were much larger than those of neat PLA and binary TPVs, resulting in a definite improvement in the toughness index. D80/10/10 obtained the highest value of 2.43, nearly 13 times than that of neat PLA. For all of the PLA/NR-PMMA/NR ternary TPVs (D80/5/15, D80/10/10 and D80/15/5), both the crack initiation energy and the crack propagation energy were considered as a significant contribution to the impact fracture toughness of the materials.39,40 The SEM images of the impact fracture surfaces in crack-initiation zone and crack propagation zone are presented in Figure S5 and Figure S6, respectively.

Figure 9. SEM micrographs of impact fractured surfaces nearby the unbroken section: (a) D80/5/15; (b) D80/10/10 and (c) D80/15/5; (d) photographs of the unbroken impact samples The morphologies of impact fracture surfaces near the unbroken section of PLA/NR-PMMA/NR ternary TPVs were also observed by SEM. As seen from Figure 9, all of the pictures showed obvious plastic deformations and stretched elastic microfibrils with high orientation. Note that the root parts of the stretched rubber microfibrils were firmly bonded with the PLA matrix without any viewable rubber cavitations, which confirmed that the rubber phase owned a specific “network-like” structure rather than dispersed particles. This partial continuous rubber phase had a super toughening effect than the conventional rubber particles on PLA.31,34 During the fracturing process, the crack energy was absorbed and dissipated by the formation of soft elastic microfibrils.

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Figure 10. Schematic of the possible toughening mechanism for continuous “soft-hard” core-shell rubber phase in PLA Based on the above study, we proposed a possible toughening mechanism for this PLA/NR-PMMA/NR ternary TPV system, which is illustrated in Figure 10. When the samples were subjected to external forces, stress concentration usually occurred in the vicinity of the rubber phase due to the difference of elastic modulus between the rubber phase and plastic phase. Thus, the effectiveness of a selected rubber toughened plastics system, in general, was mainly determined by the rubber phase and interfacial compatibility between plastic and rubber.38,41,42 In this PLA/NR-PMMA/NR ternary TPV system, the “soft” NR phase facilitated stress-transferring from the rubber phase to multiple directions, which could effectively prevent the stress concentration. As for NR-PMMA, it had a stronger interface with PLA than NR. However, the NR-PMMA phase with relative higher “rigidity” (compared with NR) was not easy to be deformed to transfer impact energies. As a result, numerous shear bands were initiated at the phase interface. Differing to the D80/0/20 and D80/20/0, the “soft-hard” core-shell rubber phase in PLA/NR-PMMA/NR ternary TPVs had a synergistic effect on energy absorption and stress dispersion.5 The “soft” NR core provided flexibility and the “hard” NR-PMMA shell imparted impact resistance, as well as sufficient plastic/rubber interfacial adhesion. When the external force was applied on the NR-PMMA shell, stress could be delivered effectively from the NR-PMMA shell layer to the NR core, and then was multi-directionally transferred to NR-PMMA shell on the other side, avoiding the stress concentration. Finally, the outer NR-PMMA phase withstood the diffused stress effectively and triggered much shear yielding in the PLA phase. So, large amounts of impact energies were dissipated during the impacting process, which resulted in tremendously enhanced impact strength for the PLA/NR-PMMA/NR ternary TPVs. Of course, the final structure of core-shell rubber phase was related to its nature strength and the compatibility with PLA, which was depended on the NR-PMMA/NR ratios. This is also the reason that an excellent tensile and flexural strength could be maintained in the mechanical performance. CONCLUSIONS In this work, bio-based PLA/NR-PMMA/NR ternary TPVs with balanced stiffness and toughness were prepared. With 10 wt % of NR and NR-PMMA, D80/10/10 displayed an enhanced yield stress of 41.7 MPa (only 38% loss compared to neat PLA) and a super impact strength of 91.3 kJ/m2 (nearly 32 times that of the neat PLA). FT-IR results confirmed that the in-situ compatibilization between PLA and rubber phases was occurred during DCP-induced dynamic vulcanization. The NR-PMMA preferred to locate between the PLA and NR phases, forming a “soft-hard” core-shell local continuous rubber phase in the PLA/NR-PMMA/NR ternary TPVs. 13



The crystallization behavior changed little when NR-PMMA/NR ratio>10/10 (w/w), which further confirmed the formation of core-shell structure. The “soft-hard” core-shell continuous rubber phase structure provided considerable energy dissipation during the impact process. At the same time, the excellent interfacial interactions between the “hard” NR-PMMA/NR shell and PLA matrix contributed to the considerable strength of the ternary TPVs. SUPPORTING INFORMATION Swelling experiments; contact angles experiments; SEM image of cryo-fracture surface of the B80/10/10 etched by dichloromethane; SEM images of the impact fracture surfaces in crack-initiation zone and crack propagation zone. ACKNOWLEDGMENT This work was supported by the National Natural Science Foundation of China [Grant No. 21704028], the Program of Guangzhou Science Technology and Innovation Commission [Grant No. 201607010103] and Program of Guangdong Provincial Department of Science and Technology [Grant No. 2016A010103004], the Natural Science Foundation of Guangxi Province (2015GXNSFBA139237), the Project Sponsored by the Scientific Research Foundation of GuangXi University (Grant No. XTZ140787), and the Fundamental Research Funds for the Central Universities.

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Synopsis: A “soft-hard” core-shell continuous rubber structure gives the bio-based PLA/NR-PMMA/NR TPVs a balanced stiffness/toughness performance.

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NR Ternary Thermoplastic Vulcanizates

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