Fully Biobased and Supertough Polylactide-Based Thermoplastic

Oct 7, 2014 - Ashish Kumar , T. Venkatappa Rao , S. Ray Chowdhury , S.V.S. Ramana Reddy. Reactive and ... Selectively cross-linked poly (lactide)/ethy...
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Fully Biobased and Supertough Polylactide-Based Thermoplastic Vulcanizates Fabricated by Peroxide-Induced Dynamic Vulcanization and Interfacial Compatibilization Guang-Chen Liu, Yi-Song He, Jian-Bing Zeng,* Qiu-Tong Li, and Yu-Zhong Wang* Center for Degradable and Flame-Retardant Polymeric Materials, College of Chemistry, State Key Laboratory of Polymer Materials Engineering, National Engineering Laboratory of Eco-Friendly Polymeric Materials (Sichuan), Sichuan University, Chengdu 610064, China ABSTRACT: A fully biobased and supertough thermoplastic vulcanizate (TPV) consisting of polylactide (PLA) and a biobased vulcanized unsaturated aliphatic polyester elastomer (UPE) was fabricated via peroxide-induced dynamic vulcanization. Interfacial compatibilization between PLA and UPE took place during dynamic vulcanization, which was confirmed by gel measurement and NMR analysis. After vulcanization, the TPV exhibited a quasi cocontinuous morphology with vulcanized UPE compactly dispersed in PLA matrix, which was different from the pristine PLA/UPE blend, exhibiting typically phase-separated morphology with unvulcanized UPE droplets discretely dispersed in matrix. The TPV showed significantly improved tensile and impact toughness with values up to about 99.3 MJ/m3 and 586.6 J/m, respectively, compared to those of 3.2 MJ/m3 and 16.8 J/m for neat PLA, respectively. The toughening mechanisms under tensile and impact tests were investigated and deduced as massive shear yielding of the PLA matrix triggered by internal cavitation of VUPE. The fully biobased supertough PLA vulcanizate could serve as a promising alternative to traditional commodity plastics.

1. INTRODUCTION Since toughness is generally one of the most important properties for structural polymer-based materials, toughening modification of brittle plastics such as polyvinyl chloride, polystyrene, polyamide, and polypropylene has been a polymer science issue for decades. As a classical way, rubber toughening, i.e., incorporation of rubbery elastomers into polymers to prepare rubber-toughened copolymers or blends, has led to varieties of commercially high-performance polymer-based materials. Meanwhile, toughening mechanisms have also been studied intensively, and a relatively thorough framework of toughening theories including multicrazes, shear yielding, craze−shear bands, cavitations and the percolation theory, etc. has been developed.1−4 However, it does not mean that simple incorporation of rubbery phase can always achieve a significant improvement in toughness, since a sharp brittletough transition at ambient temperature can be reached only when those material morphological factors (rubber particle size, volume concentration, interfacial adhesion and so on) are superior to the critical values.1−4 Therefore, controlling the final morphology becomes crucial to achieve the optimum toughening efficiency. Thermoplastic vulcanizate (TPV) is known as a particular thermoplastic elastomer with a relatively high amount of crosslinked rubber finely dispersed in the thermoplastic matrix. TPV usually shows excellent mechanical performances and processability. In fact, TPVs with thermoplastic phase as the main component have also been reported as ultratough materials.5−9 © XXXX American Chemical Society

TPVs can be fabricated via a dynamic vulcanization method which involves dynamical cross-linking of the rubber in the thermoplastic melts under high shearing dispersion. An ideal morphology with the vulcanized rubbery phase well dispersed in the plastic matrix may form during dynamic vulcanization, which is the key factor for the TPV to show excellent properties. As a biobased polymer, polylactide (PLA) has received considerable attention in both scientific research and industrial applications.10−12 PLA shows a lot of advantages such as excellent biocompatibility, biodegradability, renewability, and high mechanical strength, which makes it very attractive in the application of biomaterials and eco-friendly materials.13−15 However, the extensive application of PLA in replacing traditional commodity plastics has been hindered by its drawback of lacking toughness with small elongation at break and impact strength. Therefore, many methods have been employed to improve the toughness of PLA. Polymer blending has been proved to be the most economic and efficient way compared with plasticization and copolymerization.16−19 The most important concerns on toughening of PLA by blending are the origin and environmental compatibility of the blend component. In this sense, some biobased and/or biodegradable ductile polymers have already been used to Received: August 27, 2014 Revised: October 1, 2014

A

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2. EXPERIMENTAL SECTION

blend with PLA. For example, Zhang et al. reported a ternary blends of PLA, poly(3-hydroxybutyrate-co-hydroxyvalerate) (PHBV), and poly(butylene succinate) (PBS) with balanced properties.20 Ojijo et al. investigated the effect of in situ compatibilization of triphenyl phosphite on toughening PLA by poly(butylene succinate-co-adipate) (PLA/PBSA) blends.21 However, the improvement in impact strength seems limited even with a combination of compatibilizer or through reactive blending to obtain a better dispersion. Blending with rubbery polymers is an efficient way to improve impact strength of PLA, especially when interfacial compatibilization between PLA and the rubbery polymers took place. Oyama prepared a supertough PLLA blend with poly(ethylene-glycidyl methacrylate) (EGMA) through reactive blending.22 Dong et al. reported PLA alloys with excellent impact strength through reactive blending of PLA and ethyleneco-acrylic ester-co-glycidyl methacrylate (E-AE-GMA) rubber. The reaction between epoxy group of the rubber and the terminal group of PLA improved the compatibiliztion between the dispersed rubbery phase and the PLA matrix.23 Liu et al. fabricated a supertoughened PLA ternary blend containing ethylene/n-butyl acrylate/glycidyl methacrylate (EBA-GMA) terpolymer and zinc ionomer of ethylene/methacrylic acid (EMAA-Zn) through simultaneous dynamic vulcanization and interfacial compatibilization.24,25 Recently, we also prepared supertoughened PLA blends containing an in situ formed crosslinked polyurethane (CPU); interfacial compatibilization also occurred between PLA and the CPU through terminal group reaction.26,27 In the latter two systems, the rubbery phases are vulcanized during fabrication. The obtained materials are as a matter of fact thermoplastic-rich TPVs. It is worth noting that the above-mentioned rubbery polymers were not obtained from renewable resources; the use of those materials reduced the biobased attribute of the resulting PLA based materials. In contrast, natural rubber (NR), a biobased material, can be used to toughen PLA without compromising biobased attribute. Due to the existence of C C double bond, dynamic vulcanization of NR in PLA matrix can be induced by free radical to obtain TPV.28,29 However, the improvement in impact toughness seemed unsatisfactory. In addition, free radical reaction has also been employed to improve compatibility between PLA and some other biodegradable polymers such as poly(ε-caprolactone) (PCL),30 poly(butylene succinate) (PBS),31 and poly(butylene adipate-co-terephthalate) (PBAT).32 In this paper, we report the fabrication, morphology, compatibilization, and mechanical properties of a fully biobased TPV consisting of PLA and an unsaturated aliphatic polyester elastomer (UPE), which was synthesized from biobased monomers. Dynamic vulcanization of UPE occurred via the cross-linking of the double bonds, and interfacial compatibilization also took place through reaction of PLA with UPE in the presence of dicumyl peroxide (DCP) as an initiator. Both the tensile and the impact toughness of PLA were improved dramatically by the formation of the TPV with improved compatibiliztion and the well dispersed phase morphology. The employment of such a biobased elastomer would not affect the biobased attribute of the PLA-based materials. With significantly improved toughness, the fully biobased thermoplastic vulcanizates can find extensive application as alternatives to some conventional petroleum-based commodity plastics.

2.1. Materials. PLA 4032D was procured from NatureWorks and used without purification. The molecular weight (Mn) and dispersity (Đ) of PLA were 130 kg/mol and 2.87, respectively, obtained from the size exclusion chromatography (SEC). Itaconic acid (IA) (purity 99.0%) was received from Changzhou Xinhua Active Material Institute (Changzhou China). 1,3-Propanediol (PDO) (purity 99.0%) was purchased from Tianjin Kemiou Chemical Reagent Co., Ltd. (Tianjin China). 1,4-Butanediol (BDO, AR grade), sebacic acid (SeA, AR grade), tetraisopropyl titanate (TIPT), hydroquinone monomethyl ether (MEHQ) and DCP were obtained from Kelong Chemical Factory (Chengdu, China) and used as received. All the other chemicals with reagent grade were used without any purification. 2.2. Synthesis of the UPE. The UPE was synthesized through a two-step polycondensation procedure: esterification and subsequent polycondensation. The detailed process can be seen in studies reported by Kang et al.33,34 Briefly, BDO (0.44 mol), PDO (0.66 mol), IA (0.1 mol), SeA (0.9 mol), and the inhibitor MEHQ (0.05 wt % of total reactants) were added into a 500 mL three-necked flask equipped with a Dean−Stark water separator, a mechanical stirrer, and an inlet of dry nitrogen. The esterification reaction was carried out at 175 °C for 4h. Then, the catalyst TIPT (0.1 wt % of total amount of the reactants) was introduced and the condensation polymerization was implemented at 230 °C under 70 Pa vacuum for about 4 h until the Weisenberg effect was detected. The obtained UPE was used without any purification. The synthetic route is shown in Scheme 1.

Scheme 1. Schematic Route for Synthesis of the Unsaturated Aliphatic Polyester Elastomer

2.3. Preparation of TPV Consisting of PLA and UPE. Both PLA and UPE were vacuum-dried at 60 °C for 12 h prior to use. All the samples were prepared with a Haake Rheometer (Haake Rheomix 600, Germany) at 180 °C and 50 rpm for about 10 min. Typically, PLA and UPE with a weight ratio of 80:20 were first premixed in the chamber until the torque became stable (ca. ∼4 min), then a certain amount of DCP was added, and the dynamic vulcanization was finished until the melt torque leveled off. Subsequently, the samples were removed from the cavity of internal mixer and cooled to room temperature. A series of TPVs with the same PLA/UPE feed ratio but different amounts of DCP varying from 0 to 0.01, 0.03, 0.05, 0.1, and 0.2 phr were prepared to study the effect of DCP content on the morphology and properties of the final products. For brevity, those prepared blends were denoted as PLA/UPE and TPV-x, successively, where PLA/UPE represented the pristine binary blend without DCP and x indicated the amount of DCP in phr in TPVs. For example, TPV-0.1 represented the TPV prepared from 80 wt % of PLA, 20 wt % of UPE and 0.1 phr DCP. For property comparison, neat PLA and PLA with 0.1 phr DCP, namely, PLA-0.1, were also processed with the same conditions. Standard tensile (ASTM D638, type III) and Notched Izod impact (ASTM D256) testing specimens were prepared using a injection molding apparatus (Ray-Ran, UK) with a nozzle temperature of 220 B

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°C and a molding temperature of 60 °C. The samples were vacuumdried at 60 °C for 8 h prior to injection molding. 2.4. Proton Nuclear Magnetic Resonance (1H NMR) Spectroscopy. The chemical structure of UPE was characterized by 1H NMR (Bruker AC-P 400 MHz spectrometer) at room temperature with CDCl3 and tetramethylsilane (TMS) as a solvent and an internal chemical shift standard, respectively. In order to prove the interfacial compatibilization of PLA and UPE through chemical linkage, PLA was extracted from the TPV and characterized with 1H NMR. The measurement was also conducted at room temperature with CDCl3 and TMS as a solvent and an internal chemical shift standard, respectively. 2.5. Size Exclusion Chromatography. The molar mass (Mn) and dispersity (Đ) of the polymers were characterized by a Waters GPC, which was equipped with a 1515 pump, a Waters model 717 autosampler, and a 2414 refractive index detector. Chloroform and monodisperse polystyrene were used as the eluent and standard, respectively. The flow rate of eluent and the concentration of sample were 1.0 mL/min and 0.25 mg/mL, respectively. The experiments were carried out at 35 °C. 2.6. Gel Fraction Measurement. The injection molded samples with predetermined weight (w1) were dissolved in chloroform for 1 week and then stirred and refluxed at 70 °C for 2 days to dissolve all the unvulcanlized parts thoroughly. The suspensions were then centrifuged twice at 8000 rpm to enable the vulcanlized parts to gather at the bottom centrifuge tube. The upper layer clear solution was poured out and added with excess methanol to precipitate. The precipitates were used for NMR analysis to study the interfacial compatibilization as mentioned in Section 2.4. The obtained sediments were washed with chloroform three times and then vacuum-dried at 60 °C to a constant weight (w2). The gel fractions (Gf) were calculated by the following equation:

w Gf = 2 × 100% w1

2.10. Dynamic Mechanical Analysis (DMA). Thermo-mechanical properties of the flat sheets obtained by the compression-molding were tested in the solid state with a dynamic mechanical analyzer (DMA Q800, TA Instruments, USA) using a tensile mode. Tests were performed from −70 to 130 °C at a heating rate of 3 °C/min and an oscillation frequency of 1 Hz. 2.11. Morphological Analysis. Scanning electron microscopy (SEM; XL-30s FEG, Philips, Holland) was used to observe the morphology of the fractured surfaces of specimens. Both the impactfractured and cryo-fractured surfaces of injection-molded specimens were observed with an accelerating voltage of 5 kV. Before SEM characterization, the fractured surface was sputtered with a layer of gold. Meanwhile, the local deformations of selected tensile bars during drawing were also observed by SEM over the inside neck zones, which were longitudinally at different tensile stages. The substructural morphology of dispersed phases was studied by transmission electron microscopy (TEM, Tecnai G2F20 S-TWIN electron microscope, FEI, Holland) using stained ultrathin sections at an accelerated voltage of 200 kV. Ultrathin sections of ca. 70−80 nm in thickness from the plane perpendicular to the injection flow direction were sliced using a RMC cryo-ultramicrotome equipped with a diamond knife and mounted on Formvar-coated 200-mesh nickel grids. Osmium tetroxide (OsO4) vapor was used to selectively stain dispersed UPE or vulcanized UPE (VUPE) phases for 25 min. Analysis software ImageJ was used for the analysis of TEM images to estimate the UPE or VUPE particles size and their distribution within PLA matrix. At least 100 particles from four independent TEM images were analyzed to calculate weight-average particle diameter (dw) using the following equation: dw =

ΔΗ m − ΔΗc × 100% wf ΔHm°

(3)

where di is the particle diameter and ni is the number of particles having the particle diameter di. The cross-sectional area (Ai) of each individual particle (i) was measured and converted into an equivalent diameter of a sphere by the equation di = (4Ai /π)0.5.21,22 Additionally, particles whose sizes were too small to be properly measured at the magnification chosen were neglected. To further observe the inner deformation mechanism during impacting measurement, the impact-fractured specimen of TPV-0.1 was selected to cryogenically microtome from the stress-whitening zone immediately underneath the center section of fresh Izod impact fracture surfaces using a diamond knife. The morphology was observed with a Tecnai G2F20 S-TWIN TEM at 200 kV.

(1)

2.7. Wide-Angle X-ray Diffraction (WAXD). WAXD patterns of the extracted gel sediments were recorded with an X-ray diffractometer (Philips X’Pert X-ray diffractometer) with Cu Kα radiation. The experiment was operated at room temperature with a scan rate of 2°/ min in the angular region (2θ) of 5−40°. The gel sediments were annealed at 110 °C for 12 h to enable sufficient crystallization of PLA component prior to WAXD measurement. As a control, WAXD pattern of neat PLA was recorded with the same procedure. 2.8. Mechanical Property Measurements. All samples were preconditioned in 50% relative humidity at 23 °C for 72 h before testing. Tensile strength, elongation at break, and Young’s modulus of the samples were measured on an Instron Universal Testing Machine (Model 5567) at a crosshead speed of 5 mm/min at room temperature. At least five specimens were tested for each sample, and the averaged result was reported. Notched Izod impact testing was performed on a Sansi ZBC-50 (Shenzhen, China) impact tester at room temperature. The injection molded sample bars were made to notched specimens with 2 mm depth and 45° angle. A minimum of five samples were tested for each material, and the averaged results were reported. 2.9. Differential Scanning Calorimetry (DSC). A DSC experiment was conducted for the injection-molded samples using a TA DSC Q200 Instrument under nitrogen atmosphere. The samples of ∼5 mg ware placed into alumina crucibles and heated from 0 to 190 °C at a heating rate of 10 °C/min to measure the intrinsic degree of crystallinity (xc) of PLA matrix in the injection-molded samples according to

χc =

∑ nidi2 ∑ nidi

3. RESULTS AND DISCUSSION 3.1. Preparation of UPE. UPE was prepared by condensation polymerization of four aliphatic monomers: PDO, BDO, SeA, and IA, as shown in Scheme 1. The detailed procedure was mentioned in Section 2.2. When monomer feed molar ratio of PDO/BDO/IA/SeA was around 6/4/1/9, the obtained UPE was reported to possess preferable elastic properties.33,34 The UPE with good elasticity is anticipated to have good toughening efficiency for PLA, therefore this monomer feed ratio was employed in this study. The chemical structure of the UPE was characterized by 1H NMR. Figure 1 shows the 1H NMR spectrum of UPE. It can be seen that the resonance signals of all the characteristic protons of UPE can be found in the spectrum. Particularly, the signals of the methylene protons of the double bonds −C(CH2)− were observed at chemical shifts of 6.35 and 5.74 ppm, indicating that the pendent double bonds remained even after going through high temperature polycondensation, which was attributed to the inhibition of MEHQ. The monomeric unit molar ratios in the synthesized UPE were calculated through the ratio of integral areas of the corresponding peaks. The monomeric unit molar

(2)

where ΔHm and ΔHc are the enthalpies of melting and cold crystallization during the first heating, respectively; ΔHοm is the enthalpy assuming 100% crystalline PLA homopolymers (93.7 J/g),11 and wf is the weight fraction of PLA component in the blends. C

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Figure 2. Melt torque versus reactive blending time for neat PLA, PLA-0.1, pristine PLA/UPE blend, and the TPVs.

melt torques of preparing TPVs and processing PLA-0.1 included two steps. In the first step, a decrease in torque was observed, which was caused by the melting of PLA and UPE. After the addition of DCP, the melt torque increased gradually in the initial stage of the second step, which was ascribed to the dynamic vulcanization of UPE initiated by the free radicals formed by thermal decomposition of DCP. After going through a maximum, the torque decreased somewhat until it reached a new more or less platform at the end of mixing, implying the full vulcanization of UPE and the thermal degradation of the final TPV.36 For PLA-0.1, a limited increase in torque was also observed, which could be attributed to the branching or slight cross-linking of PLA in the presence of free radicals.37 Meanwhile, in the case of pristine PLA/UPE blend, no increase in melt torque could be observed due to the absence of free radical without adding DCP. It is worth noting that the melt torque of PLA/UPE was lower than those of neat PLA and PLA-0.1 in the first stage, which indicated that the addition of UPE could reduce the melt viscosity of PLA, suggesting that PLA had higher melt viscosity than UPE. The final melt torques of TPVs increased gradually with the content of DCP, which is reasonable since more DCP could give more free radicals and thus increase the extent of dynamic vulcanization of UPE and the excess free radicals could also increase the reaction probability of PLA with UPE, which could increase the melt torques and was preferred for interfacial compatiblization between the two polymers,30−32,38,39 as discussed in the next section. 3.3. Compatibility and Phase Morphology. 3.3.1. Thermo-mechanical Properties. DMA was used to study the thermo-mechanical properties of UPE, neat PLA, pristine PLA/ UPE blend, and the TPVs. Figure 3 presents the temperature dependence of storage modulus and tan delta of the samples. The storage modulus plots (Figure 3a) of both neat PLA and UPE showed one slope, corresponding to the α-relaxation, i.e., glass transition of neat PLA and UPE, respectively. The corresponding relaxation peak can be observed on the tan delta plots (Figure 3b). The peak temperature was defined as the glass transition temperature (Tg) and the Tg values of neat PLA and UPE were 82.0 and −34.7 °C, respectively. Unlike UPE, the storage modulus of neat PLA increased after undergoing αrelaxation, which was attributed to the cold crystallization of PLA.40 In the cases of pristine PLA/UPE blend and the TPVs, they all exhibited two slopes in the storage modulus at around −34.7 and 82.0 °C, corresponding to the α-relaxations of UPE

Figure 1. 1H NMR spectrum of the synthetic UPE.

ratio of PDO to BDO obtained from integral area ratio of peak b to peak c was 5.92:4, which was very close to the feed ratio of 6:4. The molar ratio of IA to SeA was 1:12 derived from the integral areas of peaks a plus a′ and peak e. The value was slightly lower than the feed ratio of 1:9, which might be caused by the sublimation of IA during esterification at high temperature. It is worth noting that the ratio of integral areas of peaks a and a′ to that of peak d was calculated to be 1:1, which indicates that all double bonds remained unchanged after polymerization, which is very important for the subsequent preparation of TPVs through dynamic vulcanization. The molar mass of UPE was measured by SEC. The molecular weight (Mn) and dispersity of molar mass (ĐM) of UPE were 51 kg/mol and 3.65, respectively. The differences in melt viscosity of blending components are very important in determining the morphology of the resulting blends, the melt viscosity of the dispersed phase should be lower than that of the matrix polymer to obtain a well dispersed morphology. The molecular weight of UPE was much lower than that of PLA, PLA should have higher melt viscosity than UPE, and thus blending the two polymers would give a well dispersed morphology.35 3.2. Preparation of TPV via Dynamic Vulcanization. TPV was prepared via dynamic vulcanization of UPE in PLA melting matrix with DCP as an initiator in a Haake Rheometer. For toughening modification of PLA with rubbery polymers, the feed weight ratio of PLA and the blend component was usually fixed at 80:20,19−24,27−29 which has been reported to exhibit a good balance between toughness and mechanical strength. Because if the content of rubbery polymer is too high, the blend would suffer from a large reduction in tensile strength and Young’s modulus, while if the content of rubbery polymer is too low, significant improvement in toughness especially in impact toughness is not anticipated. Therefore, weight ratio of PLA to UPE of 80:20 was used to prepare TPV in this study. Figure 2 shows the evolution of melt torques with time during dynamic vulcanization. The melt torque of neat PLA decreased significantly at the initial stage and then decreased very slightly with increasing time, which corresponded to the melting of PLA and the inevitable thermal degradation of PLA suffering from high temperature, respectively. The change in D

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VUPE) particles and loose local segmental motions, and therefore improve the toughness.41,42 It is worth noting that the Tg of the UPE component in the pristine PLA/UPE blend was about 2 °C lower than those of VUPEs in TPVs, which should be attributed to the constraining effect of cross-linked junctions on chain mobility of UPE after vulcanization. The Tg of the VUPE component seemed unaffected by the content of DCP, since they showed the same values at about −36.9 °C regardless of DCP content, which may indicate that the content of DCP would not affect the cross-linking content of the VUPE apparently, as will be confirmed by the gel fraction analysis in the following text. 3.3.2. Interfacial Compatibilization. Interfacial compatibilization is very important for the phase-separated polymer blends to show good mechanical properties. Sufficient interfacial adhesion is essential for stress transfer from one phase to the other,43 which is efficient to stop the cracks initiated at the interface from growth to catastrophic failure. In this study, the dynamic vulcanization of UPE was initiated by the free radicals formed by decomposition of DCP, which could also introduce free radicals to macromolecular chain of PLA through hydrogen abstraction.30−32,38,39 The PLA macromolecular free radicals would then reacted with CC double bonds of UPE to graft PLA onto the surface of VUPE or some unvulcanized UPE residue to form VUPE-g-PLA or UPE-g-PLA graft copolymers. In addition, branched or cross-linked PLA may also form through inter- or intramolecular reaction of PLA macromolecular free radicals. However, due to the relatively higher reactivity of double bonds with free radicals, the formation of graft copolymers should dominate. The graft copolymers located at the interface between VUPE dispersed phase and PLA matrix thus improved their compatibilities. In order to confirm the interfacial compatiblization, the content of the vulcanized part, the chemical composition of the vulcanized and unvulcanized parts, and the phase morphologies of the TPVs were investigated and compared with those of pristine PLA/ UPE blend. The contents of the vulcanized part (gel fraction) of TPVs were measured according to the procedures stated in Section 2.6. During dissolving in chloroform, clear solution was obtained for pristine PLA/UPE blend, indicating no vulcanization due to the absence of free radicals. Clear solution was also obtained for PLA with 0.1 phr of DCP, implying that no apparent cross-linking occurred, while only branching or some slight cross-linking took place as evidenced by the slight increase in the melt torque during thermal processing. Turbid suspensions were acquired for all the TPVs with addition different contents of DCP. The different solubility of PLA-0.1

Figure 3. Storage modulus (a) and tan delta (b) versus temperature for neat PLA, neat UPE, pristine PLA/UPE blend, and the TPVs.

(or VUPE) and PLA components, respectively. The increase in storage modulus could also be observed after the second relaxation due to the crystallization of PLA component. Accordingly, two relaxation peaks could be observed on the tan delta plots of pristine PLA/UPE blend and the TPVs. Interestingly, the glass transition of PLA component shifted slightly to lower temperature range compared to neat PLA, indicating some compatibility between PLA and UPE. Detailed glass transition temperature values of PLA and UPE (or VUPE) are listed in Table 1. Interestingly, Tg’s of UPE (or VUPE) components also shifted to lower temperature range compared to that of neat UPE The depression in Tg of UPE (or VUPE) component might be ascribed to the negative pressure caused by the mismatched thermal shrinkages between dispersed UPE (or VUPE) phase and PLA matrix, which would in turn cause a dilatational effect on the matrix ligament between UPE (or

Table 1. Glass Transition Temperatures and Mechanical Properties of Neat PLA, Pristine PLA/UPE Blend, and TPVs Tga

a

samples

TgUPE(or VUPE) (°C)

TgPLA (°C)

neat PLA PLA/UPE TPV-0.01 TPV-0.03 TPV-0.05 TPV-0.1 TPV-0.2

−38.61 −36.89 −36.99 −36.67 −36.82 −36.89

81.74 78.71 78.58 78.41 77.53 78.66 78.53

tensile strength (MPa) 68.1 45.3 41.2 39.2 37.7 39.9 37.5

± ± ± ± ± ± ±

1.3 1.2 2.5 2.1 2.6 3.3 2.1

Young’s modulus (MPa) 1955 1521 1445 1413 1498 1502 1501

± ± ± ± ± ± ±

28 56 60 30 92 67 40

elongation at break (%)

tensile toughness (MJ/m3)

notched Izod impact strength (J/m)

7.8 ± 0.9 302.3 ± 19.4 300.0 ± 30.5 289.5 ± 17.3 261.9 ± 26.2 257.4 ± 23.4 259.9 ± 21.8

3.2 ± 0.4 109.6 ± 5.1 99.3 ± 12.5 93.1 ± 6.5 83.9 ± 11.6 87.0 ± 9.0 84.0 ± 7.0

16.8 ± 0.7 225.1 ± 7.3 267.6 ± 23.4 484.5 ± 23.5 533.7 ± 6.0 571.2 ± 25.6 586.6 ± 26.7

Tg obtained from the DMA measurement. E

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and TPVs could prove that the reactivity of free radicals was higher with double bonds of UPE than with PLA, which indirectly indicated that once the PLA macromolecular free radicals formed, they tended to react with UPE to form grafting copolymers rather than to undergo inter- or intramolecular reaction to form self-cross-linking polymer. Figure 4 shows the gel fractions of the TPVs with addition of different contents of DCP. It can be seen that the gel fractions

Figure 4. Gel fractions for pristine PLA/UPE blend and TPVs with different contents of DCP.

Figure 5. 1H NMR spectra of neat PLA (a), unvulcanized part of TPV0.1 (b), and neat UPE (c).

of TPVs increased to a maximum ca. 30 wt % with DCP content increased to 0.1 phr. With further increasing DCP amount to 0.2 phr, the gel fraction declined to about 27 wt % for TPV-0.2. It should be noted that the drop does not mean decreased cross-linking level, since the value was higher than the content of UPE with 20 wt % in the TPVs. The theoretical gel fraction should be 20 wt % if only UPE phase was vulcanized. However, it is worth noting that the gel fractions of all the TPVs were higher than 20 wt %, which indicated that not only UPE but also some PLA entered the vulcanized part of the TPVs. Taking into account that self-cross-linking of PLA is difficult, PLA macromolecular radicals should enter the vulcanized part mainly by the way of forming grafted chains onto the surface of vulcanized UPE, namely, forming VUPE-gPLA, which would significantly improve the interfacial compatibility between the two phases. The small decrease in gel fraction of TPV-0.2 compared to that of TPV-0.1 may result from that large amount of DCP would cause much more PLA macromolecular free radicals which then increased the probability of forming soluble grafting copolymers between PLA and the not yet vulcanized UPE or forming some soluble branched PLA. To confirm the vulcanized part of TPV containing PLA component and the formation of some soluble grafting polymer, the vulcanized part was characterized by WAXD, and the soluble part was precipitated and characterized by 1H NMR. Figure 5 shows the 1H NMR spectra of neat PLA, neat UPE and the soluble part of typical TPV-0.1. Except for the two characteristic resonance signals belonging to PLA, three prominent peaks could also be found in the embedded enlarged zones of spectrum of the soluble part at the shifts of 2.15−2.45 ppm and 4.0−4.3 ppm in Figure 5b. These peaks are consistent with those in the same zones in Figure 5c, belonging to characteristic hydrogen protons of UPE, which have been labeled as “b”, “c”, and “e” in Figure 1. Considering the highly reactive free radicals were generated during the dynamic

vulcanization, the formation of soluble grafting copolymers containing PLA and UPE was reasonable. Figure 6 shows the WAXD patterns of annealed neat PLA and vulcanized parts of typical TPV-0.1. Annealed neat PLA displays an intense diffraction peak at 2θ value of 17.50°, corresponding to the (110) and/or (200) planes of α form PLA crystals.44 The characteristic diffraction peak could also be detected in the vulcanized part of TPV-0.1 although the peak

Figure 6. WAXD patterns of neat PLA (a) and the vulcanized part of TPV-0.1 (b) after annealing at 110 °C for 12 h. F

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seemed to be improved to some extent, for some sponge-like aggregates with irregular shapes could be captured. With increase in the content of DCP, the shapes of dispersed phase turned more irregular and the interfaces between VUPE and PLA became more difficult to mark, indicating improvement in interfacial adhesion between VUPE and PLA. In addition, the VUPE phase distributed much denser in TPVs than UPE in pristine PLA/UPE blend, and interconnections between VUPE phases also seemed increased, which means that VUPE phases could occupy more volumes after the addition of DCP by reducing particle size. Especially for TPV0.1 and TPV-0.2, a kind of an approximately bicontinuous structure was formed, although some dispersed VUPE phases with ellipsoidal shapes could still be observed. TEM was employed to further investigate the phase morphologies of pristine PLA/UPE blend and the TPVs. Figure 8 shows the TEM micrographs of pristine PLA/UPE

intensity was much smaller. Given that no self-cross-linking of PLA occurred during dynamic vulcanization and the free PLA was removed completely through thorough dissolution in chloroform, the specific PLA diffraction peaks in the vulcanized part could be reasonably ascribed to the formation of VUPE-gPLA, which had PLA chains grafted onto the vulcanized UPE surfaces and was located at the interface between dispersed VUPE phase and PLA matrix to improve their compatibilities. 3.3.3. Phase Morphology. The final mechanical properties of binary blends depend strongly on the phase morphologies. The phase morphologies of pristine PLA/UPE binary blend and TPVs were characterized by SEM and TEM. Figure 7

Figure 8. TEM micrographs of pristine PLA/UPE blend (a), TPV0.03 (b), TPV-0.05 (c), and TPV-0.1 (d). Figure 7. SEM micrographs for cryo-fractured surfaces of pristine PLA/UPE blend (a), TPV-0.01 (b), TPV-0.03 (c), TPV-0.05 (d), TPV-0.1 (e), and TPV-0.2 (f).

blend and some typical TPVs. The phase separation could be observed for pristine PLA/UPE blend and all the TPVs. Discrete round-like UPE droplets with obvious contour were found to disperse in PLA matrix for pristine PLA/UPE blend, implying the high surface tension and weak interfacial adhesion between UPE and PLA phases. However, after dynamic vulcanization, obvious differences were observed in both morphological features and the interfacial adhesion, which are two most important factors determining the toughness of rubber-toughened polymers.1−4 Similar to the results of SEM observation, the shape of dispersed VUPE phase become much more irregular and the contour of individual particle turns more difficult to be distinguished, indicating good interfacial compatibility between the two phases. The appearance of more clusters with interlinked dense VUPE particles revealed the highly improved continuity of the rubbery phase, which means the formation of quasi cocontinuous phase-separation morphology between VUPE and PLA.

shows the SEM images of the cryogenically fractured surfaces of pristine PLA/UPE blend and the TPVs. Neat PLA exhibits one uniform phase with a smooth surface, for brevity its SEM micrograph was not shown here. Pristine PLA/UPE blend and all the TPVs showed obvious phase-separation structures with UPE or VUPE phase dispersed in the PLA matrix with irregular shapes, implying some compatibility between the two phases. Pristine PLA/UPE blend displayed a well “sea−island” structure with discrete UPE globular “islands” embedded in PLA matrix “sea”; meanwhile, in some domains, the gaps between the surfaces of UPE and PLA phases could be easily distinguished, reflecting poor interfacial adhesion between UPE and PLA due to interfacial tension between the two phases. When 0.01 phr of DCP was added, although obvious phase separation could still be seen in TPV-0.01, the compatibility G

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Furthermore, it can be seen that the rubbery particle size reduced more or less in TPVs compared to that in the pristine binary blend. Especially for TPV-0.1, much more particles of submicron scale with irregular shapes turn up and aggregate substantially more compactly. It is known that there was an optimum size range of dispersed phase for toughening efficiency in rubber-toughened plastic systems. The optimum weight-average size (dw) of rubbery phase was in the range of 0.7−0.9 μm as been reported in the previous papers.21−24 The relationship between toughening effects and particle size could be explained as follows: when the weight-average size (dw) of dispersed phase was exactly within the range of 0.7−0.9 μm, extensive localized deformations of shear yielding or multiple crazings needed for high energy absorption upon impact tests could be promoted stably via internal cavitation, therefore highly efficient toughening effect could be obtained; While if dw deviated from the optimum value, fatal crack growth would be initiated via either unstable crazes or interfacial debondings, and propagate quickly, thus leading to poor toughness. In this study, the dw’s of dispersed phase in pristine PLA/UPE blend, TPV0.03, TPV-0.05 and TPV-0.1 were calculated as 0.77, 0.74, 0.70, and 0.65 μm, respectively. The values present a decline trends with the addition of DCP, confirming the improved compatibility of the two phases. Additionaly, all dw values are within or close to the optimum particle size range, which also suggests that high impact strengths could be obtained, as will be shown in the following sections. From the above discussion, we can conclude that the TPV consisting of PLA and vulcanized unsaturated aliphatic polyester elsatomer was prepared successfully through dynamic vulcanization of UPE in the PLA melt matrix with DCP as an initiator. During dynamic vulcanization, the hydrogen abstraction of free radicals from PLA chain give rise to PLA macromolecular free radicals, which then grafted onto the VUPE dispersed phases and some unvulcanized UPE residues through attacking the double bonds of the UPE to form grafting copolymers, which were located at the interface between VUPE and PLA phases, and significantly improved the interfacial compatibility between the two phases. Figure 9 proposes the dynamic vulcanization and interfacial compatibilization for the preparation of TPVs. 3.4. Intrinsic Crystallinity of PLA Matrix. The degree of crystallinity of PLA matrix was reported to significantly affect the toughness of toughened PLA blends.45 Therefore, the intrinsic degree of crystallinity of PLA matrix in the injection molded bars of pristine PLA/UPE blend and all TPVs were examined from the DSC heating scans, as shown in Figure 10. All the curves showed four transitions upon heating: namely, melting endothermic peak of UPE or VUPE (Tm1), glass transition of PLA matrix, cold-crystallization exothermic peak of PLA matrix (Tcc), and melting endothermic peak of PLA matrix (Tm2). Compared with neat PLA, the glass transition of the PLA phase in all blends decreased slightly, which is consistent with the DMA results, implying the partial compatibility between PLA and UPE or VUPE. What we want to focus on was the degree of crystallinity of PLA matrix of the samples. Neat PLA gained a degree of crystallinity of ∼8% during injection molding, while pristine PLA/UPE blend and the TPVs exhibited higher degrees of crystallinity with the values varied from 12.7 to 15.1% for different samples, which indicated that the incorporation of UPE or VUPE contributed to some improvements on the crystallization of PLA phase in the blends, possibly due to the enhanced chain mobility or

Figure 9. Proposed interfacial compatibilization of TPV through formation of VUPE-g-PLA during dynamic vulcanization.

Figure 10. DSC heating scans for the as-injected molded samples of neat PLA, pristine PLA/UPE blend, and the TPVs.

improved heterogeneous nucleation effect. Since the differences in the degree of crystallinity of all the samples were very small, the effect of the crystallinty on the mechanical properties of the samples can be reasonably neglected. 3.5. Mechanical Properties and Toughening Mechanism. 3.5.1. Mechanical Properties. The tensile strength, elongation at break, and Young’s modulus of neat PLA, pristine PLA/UPE blend, and all the TPVs were measured with a universal tensile tester. Figure 11 shows the stress−strain curves of the samples, and Table 1 summarizes the data for the mechanical properties. Neat PLA showed a typical curve of brittle fracture with the elongation at break, tensile strength, and Young’s modulus of 7.8%, 68.1 MPa, and 1955 MPa, respectively, while pristine PLA/UPE blend and all TPVs exhibited characteristic nature of ductile fracture with apparent yielding occurred in the stress−strain curves. It is worth noting that pristine PLA/UPE blend showed better tensile properties over all TPVs with the elongation at break, tensile strength, and Young’s modulus of 302.3%, 45.3 MPa, and 1521 MPa, which is H

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corresponding to its brittle nature. After blending with 20 wt % UPE, the impact strength was enhanced by ∼13 times to 225 J/m for the pristine PLA/UPE blend. The improvement was anticipated since the rubbery UPE was partially compatible with PLA and dispersed well in PLA matrix with the weight-average particle size of 0.77 um, which was in the optimum range in rubber-toughened PLA system. The impact strength could be further increased significantly by the formation of TPV, and the value depended strongly on the addition content of DCP, which increased to 484.5 J/m with addition of only 0.03 phr and further increased to the so-called super toughness (notched Izod impact strength is higher than 530J/m4,21,22) with the value of 533.7 J/m for TPV-0.05. The impact strength could be increased to even higher value close to 590 J/m with further increase in DCP content to 0.2 phr. The changes of phase morphology and improvement in interfacial compatibility with addition of DCP and its various contents should account for the improvement in impact strength of TPVs. 3.5.2. Toughening Mechanism under Tensile Tests. During the tensile tests, the TPVs exhibited ductile traits with stress whitening and necking through cold drawing continuously, resulting considerable tensile energy dissipation. The stress whitening was caused by the formation of crazes or “macro crazes” consisting of nano- or micro-voids and stretched polymer fibrils with high orientation. To confirm whether the stress whitening was caused by crazes or the voiding process through debonding or internal cavitation, the neck-down regions of the tensile specimens were cryo-fractured longitudinally at different stages and locations for SEM observation. Figure 12 shows the SEM micrographs of the cryo-fractured surfaces and the locations illustration. Prior to cold-drawing

Figure 11. Stress-strain behaviors of neat PLA, pristine PLA/UPE blend, and the TPVs.

reasonable since UPE is partially compatible with PLA. After dynamic vulcanization by DCP, the TPVs showed relatively poorer tensile properties. The reason for pristine PLA/UPE blend showed the highest elongation at break might be the possible plasticization effect of low-Tg flexible UPE to PLA matrix. The slightly reduced elongation at break for TPVs may be ascribed to the reduced plasticization effect due to the vulcanization of UPE. Also, the reduced tensile strength and Young’s modulus may be ascribed to the reduced rigidity of the materials through reduction in PLA matrix volume fraction accompanied by reduced particle size of VUPE. For the effect of DCP content on the tensile properties of TPVs, the elongation at break, tensile strength, and Young’s modulus changed to constant values with DCP content increased to beyond 0.05 phr, which was consistent with the change in phase morphologies of TPVs with DCP contents. The tensile toughness of the samples was calculated from the integral areas under the stress−strain curves. The values showed similar variation tendence to the elongation at break. Pristine PLA/ UPE blend showed the maximum tensile toughness of about 110 MJ/m3. The tensile toughness declined nonlinearly to 84 MJ/m3 for TPV-0.2 with addition content of DCP up to 0.2 phr. The different chemical structures of dispersed phase in pristine PLA/UPE blend and TPVs should account for the change in tensile toughness. The vulcanized structure of the dispersed phase in TPVs would somewhat obstruct internal cavitation and thus hinder the energy dissipation upon tensile load.46,47 In contrast, UPE in pristine PLA/UPE blend was not vulcanized and thus it was easier for internal cavitation to occur, inducing greater plastic deformation to dissipate tensile energy efficiently. Although the tensile properties of TPVs were somewhat poorer than pristine PLA/UPE blend, the elongation at break and tensile toughness of TPVs were improved at least by 33 and 26 times compared to those of neat PLA, respectively, indicating highly toughening efficiency in tensile toughness. Since the improvement in tensile toughness of PLA can be easily achieved through simple blending with some other flexible aliphatic polyesters, more efforts should be paid to enhance impact toughness, especially the notched Izod impact strength of PLA. Therefore, the notched impact strength of neat PLA, pristine PLA/UPE blend, and the TPVs were measured, and the results are listed in Table 1. Neat PLA showed a very small impact strength of 16.8 J/m,

Figure 12. SEM images of stretched TPV-0.1 bar at different tensile stages as schematically indicated in (d) with the double-arrow line indicating the tensile direction.

necking in Figure 12a, internal VUPE cavitations could be clearly evidenced with randomly arrayed cavities in the wider gauge at the initial stage of the tensile test; at the joint between the gauge and the narrow necking zone in Figure 12b, those random cavities become well-oriented along the tensile direction with elongated shapes by stretching. The matrix I

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around the cavities oriented to some extent, attributed to the matrix shear yielding: when it comes to the thin neck zone in Figure 12c, threadlike cavities with much higher aspect ratio and orientation could be seen; additionally, PLA matrix shearing and stretching combined with internal VUPE voiding make the dispersions and the PLA matrix difficult to discriminate, obviously indicating excellent ductility of the blends during drawing. The elongated voids with hollow cylinder-like shapes of micro scale revealed the tensile toughening mechanism of microvoids promoting shear yielding, other than crazing. Void formation could prevent the strain softening and promote the matrix shear yielding by alteration from the triaxial stress state to plane stress state, which will be discussed subsequently. 3.5.3. Deformation and Toughening Mechanism upon Impact Tests. The mechanisms of toughening various polymers with different kinds of rubbery phases have been widely studied for decades. Generally, the impact toughening effects could be directly related to the inner deformation extent of the material near the fracture sites, since high-performance blends usually display a ductile failure behavior accompanied by a large area of stress-whitening, which is caused by multicrazes or shear yielding of the matrix and implies the radical deformation of energy absorption. Meanwhile, brittle homopolymers or blends cannot deform so extensively and thus display a relatively flat failure surface without sufficient distortions. In other words, the toughening mechanism could be perceived directly through the observation of the extent of the microscopic deformation by crazes or shear yielding of the matrix, which were the two main toughening mechanisms for rubber-toughened plastics. Crazing and shear yielding as two different means of energy dissipation usually cooperate in toughening polymeric material. The shear yielding has been reported more effective in toughening modification of semicrystalline polymers. PLA, as a semicrystalline polymer, has been proven to go through much shear yielding upon impact tests in our previous study.26,27 Figure 13 shows the SEM micrographs for the impact fractured surfaces of neat PLA, pristine PLA/UPE blend, and the typical TPVs. Neat PLA showed a typical brittle fracture, and the fracture surface was smooth with few slight stress-whitening zones and fabrils. With the incorporation of UPE, pristine PLA/UPE blend and all TPVs showed much rougher surfaces, indicating a fracture transition from brittle to tough. Plastic deformation of pristine PLA/UPE blend could be evidenced by the appearance of fibril- and wrinkle-like microdomains. TPVs exhibited rougher morphologies and denser plastic deformation domains compared to pristine PLA/UPE blend, indicating improved impact strengths through absorbing more impacting energy by the formation of such rougher morphologies. It is more attractive to study the effect of DCP content on the impact fractured surfaces of TPVs. It can be seen that rougher morphologies with increased folds and faults of even intensive shear yielding could be observed with increase in DCP content within 0.05 phr, corresponding to a sharp enhancement in impact strength from 267.6 J/m for TPV-0.01 to 484.5 J/m for TPV-0.03 and 533.7 J/m for TPV-0.05. The increase in roughness of morphologies became unconspicuous with further increasing DCP content, corresponding to an indistinctive improvement in impact strength. It should be noted that matrix shear yielding that results in considerable energy absorption is just a posterior step of the fracture process, which usually contains at least two steps: formation of microvoids through particle−matrix debonding or

Figure 13. SEM images for the impact fractured surfaces of neat PLA (a), pristine PLA/UPE blend (b), TPV-0.01(c), TPV-0.03 (d), TPV0.05 (e), and TPV-0.1 (f).

internal particle cavitation and the subsequent release of high plastic constrains in the matrix through those microvoids, i.e., triggering the shear yielding. Microvoids themselves contribute little to the impact energy dissipation, but they are essential and crucial to trigger plastic deformation of the matrix, which can lead to the large energy absorption upon impact loading. In order to further confirm which kind of microvoid formation occurred during impact testing, the inner deformation of the stress-whitening zones of impact bar of the typical TPV-0.1 was observed by TEM, as shown in Figure 14. The internal VUPE cavitation was evidenced by the noticeable brightening domain inside the dispersed phase. It can be seen that nearly every VUPE particle underwent internal cavitation, while almost no

Figure 14. TEM micrographs for TPV-0.1 (a) and the stresswhitening zones of TPV-0.1after impact testing (b). J

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vulcanizates can be used as a promising alternative to traditional commodity plastics.

crazes or debonding could be observed. Additionally, the UPE phase becomes much coarser with more agglomerations emerging, which means that individual formative microvoids could extend toward the neighboring VUPE particles and also implies more overlapped matrix ligaments existing in the blends, both of which are favorable for the extensive shear yielding of the matrix around the VUPE particles. It is generally accepted that the occurrence of internal rubbery phase cavitation is largely influenced by three main factors: the interfacial adhesion, the modulus distinction, and the rubbery microdomain size.3,48−52 The internal cavitaion only happens when the interfacial adhesion is strong enough, the modulus of the dispersed phase is fairly lower than that of the matrix, and the rubbery microdomain size was at a suitable level (submicron for toughening). Otherwise, debonding microvoids, numerous unstable crazes causing premature failure, or even fatal cracks could take place. In this study the interfacial adhesion between dispersed VUPE and PLA matrix has been enhanced significantly through interfacial compatiblization, the particle size of VUPE was ∼0.7 μm, and PLA of cause showed much higher modulus at the testing termperature. That is to say the TPVs satisfied all the three conditions for internal cavitation, thus the internal cavitation triggered matrix shearing yielding was the reasonable mechanism for both tensile toughening and impact toughening.

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AUTHOR INFORMATION

Corresponding Authors

*Tel./Fax:+86-28-85410259. E-mail:[email protected]. *Tel./Fax:+86-28-85410259. E-mail:[email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by National Natural Science Foundation of China (20904034 and 51121001).

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REFERENCES

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4. CONCLUSION Fully biobased and supertough thermoplastic vulcanizates consisting of PLA and a biobased cross-linked unsaturated aliphatic polyester elastomer were prepared by dynamic vulcanization of the unsaturated polyester in the PLA melt matrix with DCP as a free radical initiator. Interfacial compatibilization between PLA and UPE took place during dynamic vulcanization through formation of VUPE-g-PLA and/ or UPE-g-PLA grafting copolymers, which were generated from the reaction of PLA macromolecular free radicals, obtained by hydrogen abstraction in the presence of DCP, toward double bonds of UPE. After dynamic vulcanization and interfacial compatiblization, the TPV exhibited quasi cocontinuous phaseseparation morphology with VUPE dispersed in PLA matrix. Meanwhile, the pristine PLA/UPE blend exhibited typically droplet-in-matrix phase-separated morphology. The sizes of dispersed VUPE phase in TPVs were somewhat smaller than that of UPE in pristine PLA/UPE blend, although the particle sizes for all samples were in the optimum range for highly toughening efficiency to PLA. Dynamic vulcanization did not seem effective for improving the tensile toughness of the blends since the pristine PLA/UPE showed higher elongation at break and tensile toughness than the TPVs; however, it was very useful to enhance the impact toughness of the blends. Compared with the pristine PLA/UPE blend, TPVs showed significant improvement in impact strength by more than 2 times. The reduced tensile toughness of TPVs compared to pristine PLA/UPE blend was ascribed to the decreased PLA matrix volume fraction and possible obstruction effect on internal cavitation after vulcanization of UPE. The tensile toughness and impact strength of TPVs increased by up to 31 and 34 times with values of 99.3 MJ/m3 and 586.6 J/m, respectively, compared to those of 3.2 MJ/m3 and 16.8 J/m for neat PLA, respectively. The toughening mechanisms under tensile and impact tests were investigated and deduced as massive shear yielding of the PLA matrix triggered by internal VUPE cavitation. The fully biobased and supertough PLA K

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