Toward Supertough and Heat-Resistant Stereocomplex-Type

Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

pubs.acs.org/Macromolecules

Toward Supertough and Heat-Resistant Stereocomplex-Type Polylactide/Elastomer Blends with Impressive Melt Stability via in Situ Formation of Graft Copolymer during One-Pot Reactive Melt Blending Shihao Deng, Hongwei Bai,* Zhenwei Liu, Qin Zhang, and Qiang Fu*

Downloaded via WEBSTER UNIV on February 8, 2019 at 18:52:42 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Chengdu 610065, P. R. China S Supporting Information *

ABSTRACT: Stereocomplexation of enantiomeric poly(L-lactide)/poly(Dlactide) (PLLA/PDLA) chains opens up a great opportunity toward sustainable PLA engineering plastic with exceptional heat resistance and durability. However, the processing and applications of stereocomplex-type PLA (SC-PLA) are significantly blocked by its inferior melt stability (i.e., the weak melt memory effect in triggering complete SC crystallization, which makes it hard to obtain exclusive formation of SC crystallites in meltprocessed products) and inherent brittleness. In this contribution, we demonstrate an unprecedented strategy to address these obstacles by one-pot reactive melt blending of the equimolar PLLA/PDLA blend with reactive poly(ethylene−methyl acrylate−glycidyl methacrylate) (E-MA-GMA) in the presence of catalyst, where both the stereocomplexation and the grafting of some PLLA/PDLA chains onto E-MA-GMA backbones take place simultaneously and competitively. Intriguingly, the E-MA-graf tPLA copolymer in situ formed can substantially improve the melt stability of SC-PLA matrix as compatibilizer, and thus highly crystalline SC-PLA/E-MA-GMA blend products with exclusive SC crystallites can be readily obtained by injection molding. Moreover, some E-MA-graf t-PLA can also strengthen the blend interface as interfacial enhancer, which gives rise to an increase in the toughening efficiency. As a result, the obtained SC-PLA/E-MA-GMA blends exhibits impressive heat resistance (the Vicat softening temperature and heat deflection temperature are as high as 201 and 174 °C, respectively) and impact toughness (the notched Izod impact strength is close to 65 kJ/m2). Notably, their comprehensive performance is superior to some commercial petroleum-derived engineering plastics. Overall, the one-pot syntheses of copolymer by in situ grafting could open up a new horizon for creating super-robust SC-PLA-based engineering plastic using industrial melt-processing technologies.

1. INTRODUCTION Over the past few decades, growing concerns over the petroleum crisis and environment pollution have stimulated the scientists to develop bioderived and biodegradable polymers as promising substitutes for traditional synthetic plastics.1,2 Polylactide (PLA), the most appealing sustainable thermoplastic completely derived from annually renewable biomass feedstocks (e.g., corn starch) and fully biodegradable in natural environment or landfills, exhibits an enormous market potential in widespread commercial applications because of its marvelous merits of desirable biocompatibility, exceptional transparency, high mechanical strength, and good processability.3−5 Up to now, PLA has been commercially used in many short-life commodity applications. However, the potential applications of PLA as a sustainable engineering plastic have been significantly blocked by its inferior heat deflection resistance (governed by the relatively low melting temperature (Tm) of homocrystallites (HCs), approximately 170−180 °C), insufficient durability (associated with the hydrolytic degradation in use), and inherent brittleness.6,7 Thus, it is increasingly urgent to overcome these blocks so as © XXXX American Chemical Society

to make PLA competitive with existing petroleum-derived engineering plastics such as polycarbonate (PC) and poly(butylene terephthalate) (PBT). Stereocomplexation between two stereoregular complementary polymers through stereoselective interaction in the crystalline state has opened up great opportunities to develop new materials with vastly enhanced physiochemical performance.8−10 PLA has two isotactic enantiomers, namely poly(Llactide) (PLLA) and poly(D-lactide). In 1987, Ikada et al.11 first discovered the stereocomplex-type PLA (SC-PLA) by blending equimolar PLLA and PDLA together. Because the SC crystallites formed from the PLLA/PDLA racemic blends possess not only stronger interchain interactions but also denser chain packing relative to the HCs, the SC-PLA exhibits many exceptional physical properties including the slightly higher mechanical strength,12,13 much better heat resistance (the heat deflection temperature can surpass 200 °C because Received: December 10, 2018 Revised: January 20, 2019

A

DOI: 10.1021/acs.macromol.8b02626 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules of the distinctively high Tm of SC crystallites, 230−240 °C),14 and substantially improved hydrolytic degradation resistance.15,16 These fascinating attributes enable the emerging applications of the SC-PLA as an engineering Bioplastic. Unfortunately, although both the high molecular weight (Mw ≥ 105 g/mol) and complete stereocomplexation of PLLA/ PDLA blends are crucial for high-performance SC-PLA products, they are mutually exclusive in the melt-processing of the blends because the high-molecular-weight (HMW) SCPLA suffers from an inferior melt stability (i.e., the poor memory effect to trigger headmost SC crystallization from the melt).14,17−20 In general, both HCs and SC crystallites are simultaneously formed when the HMW SC-PLA is completely melted and recrystallized upon subsequent cooling.14,17−21 The formation of numerous HCs in final SC-PLA products inevitably leads to a drastically deteriorated performance.14,17,22 Obviously, it is enormously challenging to achieve high-performance SC-PLA products containing 100% SC crystallites through the melt-processing of HMW PLLA/ PDLA blends.14,17,18 It has been established that the inferior melt stability of SC-PLA is associated with the phase separation between the PLA enantiomers, as evidenced by the decoupling of PLLA/PDLA chain pairs in the melt along with the drastic reduction of their interchain interactions (melt memory).20,23 To date, several strategies have been proposed to remarkably ameliorate the melt stability of SCPLA, such as stereoblock copolymerization,24 synthesis of PLAs with graft25−28 and branched architectures,29−31 selective cross-linking, 22 incorporation of nucleating agents,32−34 and addition of plasticizers or compatibilizers.35−39 For example, Pan et al.29 found that comblike HMW PLLA/PDLA blends exhibit preferential SC crystallization because of the favorable interdigitation and interactions between the enantiomeric chains. Nevertheless, the precise synthetic of PLAs with special molecular architectures is difficult, and there is scarce information available regarding their physical properties. Considering that only linear HMW PLAs have been industrially produced, recently increasing attention has been paid to the facile blending of linear PLA enantiomers with some miscible polymers, such as poly(methyl methacrylate) (PMMA)18 and poly(vinyl acetate) (PVAc).38 Unfortunately, large amounts of the compatibilizers (typically >30 wt %) are required to achieve the exclusive SC crystallization, which in turn could cause evident depression in the SC crystallinity of the PLLA/PDLA blends and resulting product performances. Therefore, it is of great significance to develop new compatibilizers capable of substantially enhancing the melt stability. The brittleness characterized by the low impact resistance and tensile toughness is the one paramount bottleneck for SCPLA to be widely used in engineering applications. Meltblending of PLA with elastomers has been frequently utilized as the most cost-effective way to considerably enhance its toughness.40−42 Unfortunately, the blending usually gives rise to a dissatisfactory toughening effect because most of these blends are incompatible, and then the weak interface makes the stress unable to be efficiently transferred from the PLA matrix to the dispersed elastomer particles during impact/ tensile deformation of the phase-separated blends.43,44 In this regard, some premade or in situ formed block/graft copolymers at the interface are required to compatibilize the blends.45−51 For instance, Oyama45 reported a supertough PLA material by reactive melt-blending of PLLA with

poly(ethylene−glycidyl methacrylate) (E-GMA). The interfacial reaction between the terminal groups of PLLA and the epoxide groups of E-GMA can lead to the in situ formation of “single comblike” graft copolymers and ultimately contributes to the superb toughness. Some “double comblike” graft copolymers with symmetric molecular structure have also been designed recently to further enhance the compatibilization efficiency in immiscible PLLA-based blends.52 So far, considerable endeavors have been devoted to develop supertough PLLA materials. However, there is still very limited work on the toughening of SC-PLA probably because it is extremely difficult to fabricate highly crystalline SC-PLA products with exclusive SC crystallites.14,17,18,53,54 In this contribution, we attempt to simultaneously enhance the melt stability and interface strength of SC-PLA/elastomer blends with a view to develop supertough and heat-resistant SC-PLA materials for potential engineering applications. To do this, the reactive poly(ethylene−methyl acrylate−glycidyl methacrylate) (E-MA-GMA) random copolymer was utilized as an elastomeric toughening agent for SC-PLA, and it has been incorporated into the mixture of PLLA and PDLA (50/ 50, w/w) by one-pot reactive melt-blending. Because the terminal hydroxyl groups of the PLAs can readily react with the epoxide groups of E-MA-GMA in the presence of catalyst,55 some PLLA/PDLA chains were expected to be grafted onto the E-MA-GMA chains at the blend interfaces during the blending process, thus in situ forming novel E-MAgraf t-PLA copolymer with both PLLA and PDLA side chains (Scheme 1). Meanwhile, the stereocomplexation between Scheme 1. Schematic Diagram Illustrating the in Situ Formation of E-MA-graf t-PLA Copolymer at the Blend Interface via Chemical Grafting of Enantiomeric PLLA/ PDLA Chains onto the E-MA-GMA Backbones during One-Pot Reactive Blending

PLLA and PDLA chains could occur competitively, depending on the blending temperature (180−220 °C). The competitive formation of the graft copolymer and SC crystallites in the blend melt has been investigated. It should be noted that some E-MA-graf t-PLA chains initially formed at the blend interfaces could be pulled into the SC-PLA matrix under the strong shearing process due to the asymmetric molecular structure.56 Formation of sufficient graft copolymer before complete stereocomplexation is very critical to improve both B


Article

Macromolecules

solvent. The operating frequency was 400 MHz. The specimens used for the 1H NMR characterization were prepared according to the separation procedure shown in Figure 1. First, the thin slices (ca. 100

the melt stability and interface strength. On one hand, the graft copolymer dispersed in the matrix could function as compatibilizer to significantly suppress the phase separation of the enantiomeric PLLA/PDLA chains and then facilitate the formation of the precursor PLLA/PDLA helical pairs upon cooling from the melt state, thus imparting the SC-PLA matrix with a substantially enhanced melt stability. On the other hand, the interface-localized graft copolymer could entangle with both the SC-PLA matrix and dispersed E-MA-GMA phase, thus strengthening the interface as interfacial enhancer due to the bridging effect. Expectedly, injection-molded SCPLA/E-MA-GMA blends with superb toughness and heat resistance have been fabricated for the first time. The outstanding comprehensive properties of the blends are also highlighted by comparing with some petroleum-derived engineering plastics. To the best of our knowledge, the role of the graft copolymer in substantially enhancing the melt stability of HMW PLLA/PDLA blends has not been reported elsewhere. We believe that the strategy proposed here could be widely applied in the design and melt-processing of other high-performance SC-PLA/elastomer blends.

2. EXPERIMENTAL SECTION

Figure 1. Flowchart showing the sample preparation procedure for 1 H NMR characterization. Note that the amorphous PLLA/PDLA/ E-MA-GMA blends were used to ensure the good solubility of PLLA and PDLA chains in some solvents.

2.1. Materials. PLLA (Mw = 1.7 × 105 g/mol, Mw/Mn = 1.7) was obtained from NatureWorks LLC, USA. PDLA (Mw = 1.3 × 105 g/ mol, Mw/Mn = 1.6) was supplied by Zhejiang Hisun Biomaterial Co. Ltd., China. The E-MA-GMA random copolymer (E/MA/GMA = 68/24/8 (wt %)) utilized as an elastomer for effectively toughening SC-PLA was purchased from Arkema Inc., France. N,N-Dimethylstearylamine (DMSA) used as the catalyst for the grafting reaction between the terminal hydroxyl groups of PLAs and epoxide groups of E-MA-GMA was provided by J&K Scientific Ltd., China. All materials used here are commercially available. Prior to reactive melt-blending, the PLLA, PDLA, and EMA-GMA pellets were vacuum-dried at 60 °C for at least 1 day to remove any moisture and avoid excessive hydrolytic degradation of PLAs. 2.2. Sample Preparation. PLLA/PDLA/E-MA-GMA blends with a trace amount (i.e., 0.3 wt %) of the DMSA catalyst were prepared by one-pot reactive melt-blending using a Haake Rheomix 600 internal mixer (Germany) operating at 200 °C and 60 rpm for 5 min. The weight ratio of PLLA/PDLA was fixed at 50/50, and the weight fraction of E-MA-GMA in the blends was varied from 0 to 20 wt %. In particular, to achieve homogeneous dispersion in the blends and resulting high catalytic efficiency for the grafting reaction, the DMSA was completely dissolved in absolute ethanol under a dry nitrogen atmosphere and then mixed with the PLLA, PDLA and EMA-GMA pellets before the blending. For convenience, the obtained reactive blends were coded as L/D/xEG-Re, where x represents the E-MA-GMA content. The unreactive PLLA/PDLA/E-MA-GMA blends without DMSA (noted as L/D/xEG-Un) were also prepared using the same method for comparison. To investigate the competition between the stereocomplexation and grafting reaction, the other two blending temperatures (i.e., 180 and 220 °C) were also used to prepare the PLLA/PDLA/E-MA-GMA blends. Standard specimens used for the property measurements were fabricated by injection molding at 240 °C using a HAAKE MiniJet II machine (Germany) with the mold temperature of 130 °C. To obtain the blends with high matrix crystallinity, the blend melts were isothermally crystallized in the hot mold for 3 min. It should be noted that the processing including reactive melt-blending at 200 °C and subsequent injection molding at 240 °C has no apparent effect on the molecular weight of PLLA and PDLA (see Figure S1). 2.3. Measurements and Characterizations. 2.3.1. 1H Nuclear Magnetic Resonance (NMR). To verify the formation of E-MA-graf tPLA copolymer during reactive blending and characterize its chain structure, 1H NMR spectra were recorded using a Bruker AVANCE III HD 400 NMR spectrometer (Switzerland) using CDCl3 as the

μm in thickness) cut from the amorphous PLLA/PDLA/E-MA-GMA blends (which were prepared by melting at 240 °C and subsequently quenching in ice water) were immersed in dioxane (where the PLAs and E-MA-graft-PLA can be selectively dissolved) for 48 h at 60 °C. Based on the 1H NMR spectra of Insoluble 1 (Figure 2d), it is proved that the Insoluble 1 is the unreacted E-MA-GMA without any trace of PLAs and E-MA-graft-PLA components. Namely, almost all PLAs and E-MA-graf t-PLA are completely dissolved in the dioxane solvent, forming Solution 1. Second, the Solution 1 was poured into methyl alcohol. After precipitation and vacuum-drying, the obtained Deposit 1 was weighed. Then one part of Deposit 1 was dissolved in CDCl3, and the 1H NMR spectrum indicates the coexistence of unreacted PLAs and EMA-g-PLA components in Deposit 1 (Figure 2e). The other part of Deposit 1 was immersed in cyclohexane capable of selectively dissolving the EMA-g-PLA at 40 °C for 48 h, forming Insoluble 2 and Solution 2. Third, the Solution 2 was poured into methyl alcohol for precipitation, and Deposit 2 was obtained after being vacuum-dried. The Insoluble 2 and Deposit 2 were dissolved in CDCl3 to gain 1H NMR spectra (Figure 2f,g). The results indicate that the components of the Insoluble 2 and Deposit 2 are unreacted PLAs and E-MA-graft-PLA, respectively. 2.3.2. Scanning Electron Microscopy (SEM). The phase morphology was observed with a FEI Inspect F SEM (USA) operating at 5 kV. The specimens were prepared by cryo-fracturing of the injection-molded bars in liquid nitrogen, and then the cryofractured surfaces were etched by cyclohexane at 40 °C to selectively remove the E-MA-GMA phase. The impact-fractured surfaces of the specimens after the notched Izod impact test were also observed with the SEM. Before SEM imaging, both the cryo-fractured and impact-fractured surfaces were coated with a thin gold layer. 2.3.3. Transmission Electron Microscopy (TEM). A FEI Tecnai G2 F20 TEM (USA) was also used to observe the phase morphology at 200 kV. Ultrathin specimens (ca. 80 nm in thickness) were prepared by cryo-ultramicrotomy from the injection-molded bars at −100 °C using a Leica UCT microtome (Germany), followed by staining with the vapor of RuO4. The E-MA-GMA domains selectively stained by RuO4 could become gray under the TEM observation. C


Article

Macromolecules

Figure 2. 1H NMR spectra for (a) pure PLA, (b) pure E-MA-GMA, (c) reactive PLLA/PDLA/15E-MA-GMA blend with 0.3 wt % catalyst, and (d−g) the components separated from the blend. WAXD crystallinity of HCs (XWAXD c,HC ) and SC crystallites (Xc,SC ) was also estimated by comparing the peak area of HC and SC diffractions with the total areas of crystalline and amorphous phases, respectively. 2.3.6. Dynamic Mechanical Analysis (DMA). DMA was conducted on a TA Q800 instrument (USA) in a single-cantilever mode with an oscillating frequency of 1 Hz and an amplitude of 10 μm. The dynamic storage modulus (G′) was recorded at a heating rate of 3 °C/min from 0 to 230 °C. 2.3.7. Vicat Softening Temperature (VST). The VST measurement was performed using a Coesfeld HDT-Vicat 40-197-100 tester (Germany) equipped with a flat-ended needle of 1 mm2 cross section at a heating rate of 2 °C/min. The measured VST is the temperature at which the needle perforates the specimen (ca. 4 mm in thickness) to the depth of 1 mm under a constant load of 10 N. At least three replicated specimens were measured for each sample, and the average values obtained are presented. 2.3.8. Heat Deflection Temperature (HDT). The HDT was measured using an HDT/V-3116 tester (China) according to the ISO 75 standard. The measurement was performed in a flatwise mode (64 mm in the span length) at a heating rate of 2 °C/min. The loading stress is 0.45 MPa. 2.3.9. Mechanical Tests. The notched Izod impact strength was evaluated using a VJ-40 impact tester unit (China), and the tensile properties were measured using a SANS universal testing instrument (China) at a constant crosshead speed of 5.0 mm/min. For each sample, the mechanical tests were performed at ambient temperature (23 °C) on six replicated specimens, and the average test values are reported.

2.3.4. Differential Scanning Calorimetry (DSC). Thermal analysis was performed using a PerkinElmer pyris-1 DSC (USA) under a dry nitrogen gas flow (40 mL/min). For the analysis of melt stability, each specimen (ca. 5 mg) was first heated from 30 to 250 °C and maintained at this temperature for 5 min to erase thermal history. It was then cooled to 30 °C at a scanning rate of 10 °C/min and heated to 250 °C again at a rate of 10 °C/min to examine the meltcrystallization and subsequent melting behavior. In some cases, the DSC nonstop heating−cooling−heating cycle in the temperature range 30−250 °C was repeated at least three times to fully evaluate the melt stability. To analyze the matrix melting behaviors of the blends, the specimens were heated from 30 to 250 °C at a rate of 10 °C/min. Then, by use of the melting enthalpy of HCs (ΔHm,HC) and SC crystallites (ΔHm,SC) obtained from the DSC heating runs, the DSC DSC crystallinity of HCs (XDSC c,HC) and SC crystallites (Xc,SC ) in the blends can be calculated according to the following equations: DSC Xc,HC =

DSC Xc,SC =

ΔHm,HC 0 wf ΔHm,HC

(1)

ΔHm,SC 0 wf ΔHm,SC

ΔH0m,HCand

ΔH0m,SC

(2) 8

8

(taken as 93 J/g and 142 J/g ) are the where melting enthalpy for an infinitely large homocrystallite and SC crystallite, respectively; wf is the weight percent of PLAs. The relative fraction of SC crystallites (f SC) in PLLA/PDLA matrix was evaluated by

fSC =

DSC Xc,SC DSC Xc,SC

DSC + Xc,HC

3. RESULTS AND DISCUSSION 3.1. Formation of Copolymer by in Situ Grafting Reaction. To verify the formation of E-MA-graf t-PLA copolymer during melt blending of PLLA/PDLA/E-MAGMA blends in the presence of DMSA catalyst, 1H NMR measurement was performed after completely separating the components of the reactive blends. Figure 2 shows the 1H

(3)

2.3.5. Wide-Angle X-ray Diffraction (WAXD). The crystal structure was examined using an X’Pert pro MPD X-ray diffractometer (Holland) equipped with a Cu Kα source operating at 40 kV and 40 mA. The WAXD spectra were collected in the 2θ range of 5°−40° at a scanning rate of 5°/min. The WAXD D


Article

Macromolecules

Table 1. Chemical Characteristics for the Formation of E-MA-graf t-PLA in PLLA/PDLA/15E-MA-GMA Blends Prepared at Different Temperatures temp (°C)

M0 (g)

M1 (g)

RDeposit1

FNMR (%)

Fweight (%)

RDeposit2

nPLA/nEG

180 200 220

10.0025 10.0043 10.0076

1.4657 1.3772 1.3728

462.39 120.34 115.92

2.1 7.9 8.2

2.3 8.2 8.6

10.59 38.61 39.11

2.1 7.7 7.8

Figure 3. Plots of (a) the content of SC crystallites and (b) weight fraction of reacted E-MA-GMA as a function of blending time during reactive blending of PLLA/PDLA/15E-MA-GMA blends at different temperatures.

NMR spectra of pure PLA, pure E-MA-GMA, and the components separated from the PLLA/PDLA/15E-MA-GMA blend. The characteristic resonance peaks at around 1.7 and 5.2 ppm are assigned to the methyl and methine protons of PLAs unit (Figure 2a), respectively, while the characteristic peak at around 3.7 ppm is associated with the methyl proton of MA unit in E-MA-GMA (Figure 2b). The 1H NMR spectra presented in Figure 2c−g confirm that although unreacted EMA-GMA (Insoluble 1) and PLAs (Insoluble 2) components are separated from the PLLA/PDLA/15E-MA-GMA blend, the remaining Deposit 2 contains not only E-MA unit but also the PLAs unit. Because the cyclohexane is not a good solvent for PLAs, the presence of the PLAs unit in the Deposit 2 (obtained from Solution 2) clearly indicates that the in situ formation of E-MA-graf t-PLA via grafting of some PLAs chains onto the E-MA-GMA backbone. Because the E-MA-GMA is utilized as elastomer to effectively toughen SC-PLA, only small amounts of E-MAGMA chains are expected to be involved in the grafting reaction during reactive melt-blending of PLLA/PDLA/EMA-GMA blends. In this case, it is necessary to measure the amount of E-MA-GMA reacted during the blending. One simple method to calculate the weight fraction of the reacted E-MA-GMA (Fweight) is to weight according to the expression Fweight =

FNMR =

(5)

where RDeposit1 is the peak area ratio of methine of PLAs units to methyl of MA units, M0,PLA is the molar mass of PLAs repeating unit (72 g/mol), and M0,EG is the molar mass of EMA-GMA repeating unit (360.8 g/mol). The results obtained from the two calculation methods are listed in Table 1. It is interesting to find that the Fweight value (8.2 wt %) is consistent with the FNMR value (7.9 wt %), implying that about 8% of E-MA-GMA chains are involved in the grafting reaction during reactive blending of the PLLA/PDLA/15EMA-GMA blend. The chain structure of in situ formed E-MA-graf t-PLA copolymer is also characterized by using the 1H NMR spectrum of Deposit 2 (Figure 2g). The mole ratio between PLAs side chains and E-MA-GMA backbone in the E-MAgraf t-PLA can be calculated by M 0,PLA MEG nPLA = 3RDeposit2 nEG M 0,EG MPLA

(6)

where MPLA and MEG are the average molecular weights of PLAs (1.5 × 105 g/mol) and E-MA-GMA (5 × 104 g/mol), respectively. The calculated result (Table 1) indicates that the average number of the long PLLA/PDLA chains grafted onto each short E-MA-GMA backbone is ∼7.7. Namely, the E-MAgraf t-PLA copolymer is composed of 7−8 PLAs side chains along the E-MA-GMA backbone. In this case, the average molecular weight of the E-MA-graft-PLA is about 1100−1250 kg/mol. Besides, assuming 8 wt % E-MA-GMA is involved in the grafting reaction, the weight fraction of the in situ formed E-MA-graf t-PLA (Fgraft) in the PLLA/PDLA/15E-MA-GMA blend is estimated to be as high as 28.9 wt % using the following equation:

Wf,EGM 0 − M1 Wf,EGM 0

M 0,EG M 0 − M1 Wf,EGM 0 (3RDeposit1 × M 0,PLA ) + M 0,EG

(4)

where M0 and M1 are the weights of PLLA/PDLA/15E-MAGMA blend and dried Insoluble 1 (i.e., unreacted E-MAGMA), respectively; Wf,EG is the weight fraction of entire EMA-GMA in the blend. The weight fraction of the reacted EMA-GMA can also be estimated by taking use of the 1H NMR spectrum (Figure 2e), based on the calculation of the E-MAGMA fraction in Deposit 1. The relative content of PLAs over E-MA-GMA can be estimated by comparing the characteristic peak area of the methine proton of PLAs units and the methyl proton of MA units. Then, the weigh fraction of the reacted EMA-GMA (FNMR) can be calculated by

Fgraft = 0.08 × 15 × E

7.7MPLA + MEG × 100% MEG

(7)


Article

Macromolecules

Figure 4. SEM and TEM micrographs showing the phase morphology of PLLA/PDLA/15E-MA-GMA blends (a−a″) with and (b−b″) without catalyst.

3.2. Competition between SC Crystallization and Grafting Reaction during Reactive Melt-Blending. Besides the in situ grafting reaction, triggering complete SC crystallization during the one-spot reactive blending of PLLA/ PDLA/E-MA-GMA blends is of vital importance for the preparation of SC-PLA/elastomer blends through meltblending. It has been proved that the blending temperature plays a key role in regulating SC crystallization between PLLA and PDLA, and pure SC-PLA without any HCs can be easily achieved at a special temperature window between the Tms of SC crystallites and HCs.17,57 Nevertheless, the grafting reaction and SC crystallization would be competing during reactive blending because the formation of solid-state SCcrystallite particles in the blend melts could significantly suppress the grafting of PLLA/PDLA chains onto the E-MAGMA backbones. In such a case, it is necessary to facilitate the grafting of some PLLA/PDLA chains before their SC crystallization so as to obtain sufficient amounts of E-MAgraf t-PLA. Thus, the temperature dependence of the competing formation of E-MA-graf t-PLA and SC crystallites in the reactive blending has been monitored. Figure 3 shows the variations of SC content (XWAXD c,SC ) and weight fraction of reacted E-MA-GMA (FNMR) during reactive blending of the PLLA/PDLA/15E-MA-GMA blends at various temperatures. WAXD Expectedly, at a low temperature of 180 °C, the Xc,SC increases dramatically from 8.2% to 43.8% at the early stage of blending (within 45 s), and then only a small fraction of EMA-GMA chains (ca. 2.1%) can be involved in the grafting reaction to form E-MA-graf t-PLA, suggesting that formation of numerous solid-state SC crystallites could substantially suppress the grafting reaction. By contrast, when the blending temperature is increased to 200 °C, the XWAXD remains very c,SC low (ca. 10.8%) after 45 s, which provides sufficient time for more E-MA-GMA to take part in the grafting reaction (the FNMR can reach to 7.9%) because the grafting reaction mainly occurs in the melt state. It implies that depressing SC crystallization rate is favorable for the formation of numerous E-MA-graf t-PLA. However, with further increasing the blending temperature up to 220 °C, more E-MA-GMA can

immediately react with the PLAs at the early stage of blending but the SC crystallization becomes extremely difficult (no SC crystallites can be detected after blending for 90 s). Notably, the increasing in the blending temperature from 200 to 220 °C does not lead to an obvious increase in the total amount of the reacted E-MA-GMA (the maximum FNMR is ca. 8%), and also the average number of PLAs side chains in the formed EMA-graf t-PLA remains identical (ca. 8, Table 1). These results distinctly indicate that both the efficient grafting reaction and complete SC crystallization can readily proceed during reactive blending of PLLA/PDLA/E-MA-GMA blends at an optimum temperature of 200 °C. 3.3. Phase Morphology. The phase morphology of ternary PLLA/PDLA/E-MA-GMA blends and those with 0.3 wt % catalyst was observed with SEM and TEM. Figures 4a and 4b show some representative SEM micrographs of the PLLA/PDLA/15E-MA-GMA blends with and without catalyst, in which the E-MA-GMA phase was selectively etched by cyclohexane. Noticeably, the two blends exhibit the same “sea−island” morphology (i.e., spherical E-MA-GMA droplets are homogeneously distributed in the PLLA/PDLA matrix), but the average E-MA-GMA droplet size decreases evidently (from 0.5 to 0.3 μm) with the homogeneous dispersion of 0.3 wt % catalyst into the PLLA/PDLA/15E-MA-GMA blend. Because the PLAs and E-MA-GMA are immiscible, the catalyst induced grafting reaction between PLLA/PDLA and E-MA-GMA chains can only take place at the blend interfaces during reactive blending, and thus the decreased droplet size should be ascribed to the in situ formation of interfacelocalized E-MA-graf t-PLA copolymer, which is further supported by the TEM micrographs presented in Figures 4a′ and b′. In the TEM micrographs, the dark domains represent the E-MA-GMA droplets selectively strained by RuO4, the gray shells around the dark ones are the partially strained E-MA-graf t-PLA copolymer, and the white phase is the PLLA/PDLA matrix. It is interesting to find that the reactive blending induces the formation of interface-localized E-MA-graf t-PLA (as marked by the arrows in Figure 4a″), while such E-MA-graf t-PLA phase cannot be distinguished F


Article

Macromolecules

Figure 5. DSC thermograms of PLLA/PDLA/E-MA-GMA blends (a−a″) with and (b−b″) without catalyst: (a, b) during the first heating scan, (a′, b′) during the cooling scan after being completely melted at 250 °C, and (a″, b″) during the subsequent heating scan. The heating/cooling rate was 10 °C/min.

PDLA matrix during the melt-blending at 200 °C. However, after being melted at 250 °C, the unreactive blends exhibit two characteristic melting peaks in the second heating curves at around 160−180 and 200−220 °C (corresponding to the HCs and SC crystallites, respectively) (Figure 5b″). Moreover, the relative fraction of SC crystallites (f SC) is lower than 40%, implying the predominant formation of HCs upon cooling from the blend melts (Figure 5b′) due to the inferior melt stability. With regard to the slight increase in the f SC of PLLA/PDLA (50/50) blend with the incorporation of E-MAGMA, it may be ascribed to the promoting effect of the flexible E-MA-GMA on the mobility and mutual diffusion between PLLA and PDLA chains as an plasticizer.35,58 Impressively, the formation of E-MA-graf t-PLA copolymer in the reactive blends gives rise to a remarkable increase in the f SC (Figure 5a″), suggesting that the E-MA-graft-PLA can substantially enhance the melt stability of PLLA/PDLA matrix. Moreover, the f SC exhibits a strong dependence on the E-MA-GMA content, distinctly indicating that the formation of sufficient amounts of E-MA-graf t-PLA is necessary for achieving exceptional melt stability. Most importantly, the exclusive SC crystallization can be obtained in the PLLA/PDLA/E-MA-GMA blends containing relatively

from the blend without catalyst (Figure 4b″) (no E-MA-graf tPLA chains can also be detected by 1H NMR because of the low reactivity between the terminal hydroxyl groups of PLAs and epoxide groups of E-MA-GMA in the absence of catalyst). It provides a direct evidence of the formation of E-MA-graf tPLA at the interfaces. The interface-localized E-MA-graf t-PLA could suppress the coalescence of E-MA-GMA droplets and enhance the interfacial adhesion as an effective interfacial compatibilizer, ultimately contributing to the properties of the reactive blends. In addition, it should be noted that not all of the in situ formed E-MA-graf t-PLA can be stably located at the interfaces due to their highly asymmetric molecular structure, and then many E-MA-graf t-PLA chains could be pulled into the SC-PLA matrix by strong shearing forces involved in the reactive blending.56 3.4. Melt Stability. The melt stability of PLLA/PDLA/EMA-GMA blends with and without catalyst has been comparatively investigated by DSC, and some representative DSC thermograms are presented in Figure 5. As expected, for both the unreactive and reactive blends, only one melting characteristic peak of SC crystallites can be observed in the first heating curves at around 200−220 °C (Figure 5a,b), further confirming the complete SC crystallization of PLLA/ G


Article

Macromolecules

Figure 6. DSC thermograms of PLLA/PDLA/15E-MA-GMA blends (a) with and (b) without catalyst during the nonstop heating−cooling− heating cycles (the heating/cooling rate was 10 °C/min). (c) Schematic diagram illustrating the possible mechanism for the dramatic enhancement in the melt stability of SC-PLA matrix by the E-MA-graft-PLA copolymer.

Figure 7. (a, c) DSC heating curves and (b, d) WAXD patterns of the injected-molded PLLA/PDLA/E-MA-GMA blends with various amounts of E-MA-GMA.

high contents (e.g., 15−20 wt %) of E-MA-GMA, in which the concentration of the E-MA-graf t-PLA is about 28.9−38.5 wt %, as evidenced by a single characteristic melting peak of SC crystallites in the second heating curves at 200−220 °C. Obviously, the E-MA-graf t-PLA exhibits an unparalleled

enhancing effect on the melt stability but its efficiency seems very low because the exclusive SC formation are mainly resulted from the E-MA-graf t-PLA dispersed in the PLLA/ PDLA matrix. The SC crystallization temperature (Tc,SC) is ∼135 °C (Figure 5a′). H


Article

Macromolecules

Figure 8. (a) Storage modulus (G′) vs temperature curves and (b) Vicat softening temperature (VST) as well as heat deflection temperature (HDT) of highly crystalline PLLA and PLLA/PDLA15/E-MA-GMA blends with and without catalyst.

molded blends. As shown in Figure 7a,c, all blends possess the same high level of overall crystallinity (i.e., 34−41% by DSC) but diverse crystal compositions. For the unreactive PLLA/ PDLA/E-MA-GMA blends, the mixture of HCs and SC crystallites is formed in the PLLA/PDLA matrix during injection molding due to the insufficient melt stability (Figure 7c). However, with the substantial enhancement in the melt stability by E-MA-graf t-PLA, the SC crystallites are predominantly formed in the reactive PLLA/PDLA/E-MA-GMA blends (Figure 7a). In particular, SC crystallites are exclusively formed in the matrices of both PLLA/PDLA/15E-MA-GMA and PLLA/PDLA/20E-MA-GMA blends. Similar results can also be obtained from the WAXD patterns shown in Figure 7b,d. After injection molding, the unreactive blends exhibit not only strong characteristic diffractions of α-form HCs appear at 16.9°, 19.1°, and 22.5° (assigned to the (200)/ (110), (203), and (210) crystal planes, respectively) but also weak characteristic diffractions of SC crystallites at 12.1°, 20.9°, and 24.0° (belonging to the (110), (300/030), and (220) crystal planes, respectively) (Figure 7d). By contrast, in the case of the reactive blends containing abundant E-MAgraf t-PLA chains (28.9−38.5 wt %), the SC diffractions become overwhelming along with the gradual disappearance of of the HC diffractions (Figure 7b). Moreover, the XWAXD c,SC the blends can reach to 43−45%. The heat resistance of the injection-molded PLLA/PDLA/ E-MA-GMA blends was evaluated by DMA data, Vicat softening temperature (VST), and heat deflection temperature (HDT). For comparison, the injection molded PLLA with the same high crystallinity was also prepared. As shown in Figure 8a, the storage modulus of highly crystalline PLLA starts to drop considerably once the temperate approaches to ca. 151 °C (below the Tm,HC). When it comes to the unreactive PLLA/PDLA/15E-MA-GMA blend containing 42% HC crystallites and 58% HCs, the presence of many SC crystallites can slightly increase the heat distortion temperature but still has no effective resistance to heat distortion near the Tm,HC. However, for the reactive blend containing 100% SC crystallites, the considerable drop in the storage modulus can be observed only when the temperature reaches a much higher temperature of ca. 209 °C (close to the Tm,SC), indicating a markedly enhanced heat distortion resistance. The VST measurements give similar results as expected. The VSTs of the highly crystalline PLLA and unreactive PLLA/PDLA/ 15E-MA-GMA blend are only 143 and 151 °C, respectively (Figure 8b). By contrast, the VST of the reactive PLLA/ PDLA/15E-MA-GMA blend is as high as 201 °C. The HDT

The DSC cyclic experiments were also performed to further verify the exceptional melt stability of the reactive blends. As shown in Figure 6a, the crystallization in the form of SC crystallites after complete melting is perfectly reversible (the Tm,SC, ΔHm,SC, and Tc,SC keep constant in the heating− cooling−heating thermal cycles), without any HCs. To our best knowledge, this is the first paradigm to prepare HMW PLLA/PDLA/elastomer blends with inspiring melt stability via one-spot reactive blending. The role of the E-MA-graf t-PLA in enhancing the melt stability of PLLA/PDLA matrix is not yet clear, and further investigations are now in progress. The possible mechanism is schematically illustrated in Figure 6c. According to the reported packing models of helical PDLA and PLLA chains in the SC crystallites, the key to trigger complete SC crystallization upon cooling is believed to be the generation of abundant PLLA/PDLA chain pairs59 and/or chain assemblies with nonequivalent PLLA/PDLA ratio60 in the melt. For the blends without E-MA-graf t-PLA, the PLLA/ PDLA pairs and assemblies could be completely destroyed when the SC crystallites are melted and even undergo microphase separation,20,23 thus forming many PLLA- and PDLA-rich domains in the blend melts. In this case, the PLLA or PDLA chains located in these individual domains may prefer to taking part in the homocrystallization within the domains rather than the SC crystallization by overcoming mutual diffusion barrier between adjacent domains (the PLLA and PDLA chains must suffer from remarkably prolonged diffusion pathway toward the growth front of SC crystallites). As a result, both the HCs and SC crystallites are simultaneously formed. However, the E-MA-graft-PLA chains formed in the reactive blends could serve as efficient compatibilizer to prevent the PLLA and PDLA chains completely decoupling from the helical pairs and assemblies after melting of SC crystallites by facilitating their interchain interactions. The substantial depression of the microphase separation makes the PLLA and PDLA chains able to readily collaborate with each other and then assemble into the precursor PLLA/PDLA helical pairs in the melt; thus, the blends exhibit an impressive ability to exclusively form SC crystallites during melt crystallization. 3.5. Heat Resistance. To elucidate the importance of substantially enhancing melt stability in the production of high-performance SC-PLA/elastomer blends, highly crystalline PLLA/PDLA/E-MA-GMA products with and without catalyst were prepared by injection molding. Figure 7 presents the DSC heating curves and WAXD patterns of these injectedI


Article

Macromolecules of the blend can reach to 174 °C, about 53 °C higher than that of the unreactive counterpart, which is highly desirable in various engineering applications. These results vividly demonstrate that the excellent melt stability is indispensable for fabricating heat-resistance SC-PLA products by meltprocessing. 3.6. Mechanical Properties. As mentioned above, the EMA-graf t-PLA copolymer formed at the interface of the PLLA/PDLA/E-MA-GMA blends could function as interfacial enhancer to enhance the interfacial adhesion. To elucidate the influence of the enhanced interfacial adhesion on the mechanical response of these blends under impact and uniaxial tensile tests, both the notched impact toughness and tensile properties were measured. Figure 9 displays the

Figure 10. SEM micrographs of the impact fractured surfaces for PLLA/PDLA/15E-MA-GMA blends (a, a′) with and (b, b′) without catalyst.

E-MA-graf t-PLA enhanced interfacial adhesion makes the stress transferring along the interfaces more efficient, and thus much more fracture energy is dissipated during the impact fracture process. Mechanical strength and toughness are two mutually exclusive attributes of polymer materials. In most cases, efficient toughening is obtained at the cost of greatly reducing the strength and stiffness. Figure 11 depicts the typical stress−

Figure 9. Notched Izod impact strength of PLLA/PDLA/E-MAGMA blends with various amounts of E-MA-GMA.

notched Izod impact strength of PLLA/PDLA/E-MA-GMA blends with various amounts of E-MA-GMA. Obviously, the incorporation of E-MA-GMA into PLLA/PDLA matrix induces a dramatic enhancement in the impact toughness. For example, the notched Izod impact strength of PLLA/ PDLA/20E-MA-GMA blend is 45.3 kJ/m2, much higher than that (2.5 kJ/m 2 ) of the PLLA/PDLA matrix. More intriguingly, with the formation of E-MA-graf t-PLA in the blend, the toughening of dispersed E-MA-GMA particles on the PLLA/PDLA matrix becomes more efficient because the interface-localized E-MA-graf t-PLA could enhance interfacial adhesion and decrease the E-MA-GMA particle size, especially when the E-MA-GMA content reach 15 and 20 wt %. Impressively, the reactive PLLA/PDLA/20E-MA-GMA blend shows a superior notched impact toughness of as high as 64.8 kJ/m2. To get an in-depth insight into the prominent role of the E-MA-graf t-PLA enhanced interfacial adhesion in the toughening, the impact-fractured surfaces of the blends were observed with SEM and some representative micrographs are presented in Figure 10. The unreactive PLLA/PDLA/15EMA-GMA blend shows a rough fracture surface, but the significant plastic deformation does not percolate throughout the PLLA/PDLA matrix (Figure 10b,b′). One can clearly observe some regions of brittle fracture and the interfacial debonding at the interfaces (as highlighted by the arrows in Figure 10b′), However, for the reactive PLLA/PDLA/15EMA-GMA blend possessing superior impact toughness, the significant matrix plastic deformation can be seen on the whole fractured surface along with the disappearance of the interfacial debonding (Figure 10a,a′). This suggests that the

Figure 11. Tensile strain−stress curves of PLLA/PDLA/E-MA-GMA blends with and without catalyst.

strain curves of some unreactive and reactive PLLA/PDLA/EMA-GMA blends. Clearly, the PLLA/PDLA blend has high tensile strength and Young’s modulus but very low elongation at break which can be significantly enhanced by incorporating E-MA-GMA. Very intriguingly, although the tensile strength and modulus decrease greatly with increasing E-MA-GMA content, the formation of E-MA-graf t-PLA is found to give rise to an evident enhancement in the strength and elongation at break of the PLLA/PDLA/E-MA-GMA blends because the enhanced interfacial adhesion can prevent the debonding of EMA-GMA particles from the matrix at relative low stress and strain levels. Namely, the formation of E-MA-graft-PLA can simultaneously enhance the impact toughness, tensile strength, J


Article

Macromolecules

molded PLLA/PDLA/E-MA-GMA blends with much superior Vicat softening temperature (ca. 201 °C) and notched Izod impact toughness (ca. 65 kJ/m2) have been obtained for the first time due to the excusive formation of high-content SC crystallites and the great enhancement in the interface adhesion, as compared to those (ca. 151 °C and 45 kJ/m2, respectively) of the unreactive counterparts. These results reveal the importance of elastomer-graf t-PLAs copolymer in developing high-performance SC-PLA/elastomer blends by industrially meaningful melt processing, which could provide a promising avenue toward commercial-scale production of super-robust SC-PLA engineering Bioplastic.

and elongation at break. For example, the tensile strength and elongation at break of the PLLA/PDLA/15E-MA-GMA blend are increased from 39.2 MPa and 15.1% to 47.5 MPa and 26.2%. On the basis of the above results, it is clear that PLLA/ PDLA/E-MA-GMA blends with excellent heat resistance, superb toughness, and high strength can be easily prepared by one-spot reactive blending and subsequent melt-processing because the E-MA-graf t-PLA copolymer in situ formed in the blends can not only substantially enhance the matrix melt stability but also reinforce the interfaces. Figure 12 highlights

■

ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b02626. Details on the measurement of GPC, and the GPC data of PLAs before and after melt-processing (PDF)

■

AUTHOR INFORMATION

Corresponding Authors

*Tel/Fax +86 28 8546 1795; e-mail [email protected], [email protected] (H.W.B.). *Tel/Fax +86 28 8546 1795; e-mail [email protected] (Q.F.)

Figure 12. Comparison of the comprehensive performance (in terms of tensile strength, notched impact strength, and Vicat softening temperature) between PLLA/PDLA/15E-MA-GMA blend and other commercial engineering plastics (PA6,61,62 PC,63 and PBT64,65).

ORCID

Hongwei Bai: 0000-0003-4927-6422 Qiang Fu: 0000-0002-5191-3315

the superiority of the blends over the widely used petroleumderived engineering plastics (e.g., PA6, PC, and PBT) in terms of the comprehensive performance. Although the PC and PBT have superior strength, their impact toughness is much lower than that of the PLLA/PDLA/15E-MA-GMA blend.

Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work was financially supported by the National Natural Science Foundation of China (Nos. 51873129 and 51721091).

4. CONCLUSIONS In conclusion, supertough and heat-resistant PLLA/PDLA/EMA-GMA blends with excellent matrix melt stability and appropriate interfacial adhesion have been successfully prepared by one-spot reactive blending. During the blending process, not only the SC crystallization between PLLA and PDLA chains but also the grafting reaction between the PLAs and E-MA-GMA chains in the presence of catalyst occurs simultaneously and competitively. We observe that the reactive blending temperature plays a key role in controlling the competition, and the grafting becomes predominant with increasing the blending temperature from 180 to 220 °C. To be specific, the grafting of many PLAs chains onto the E-MAGMA main chains can proceed prior to their SC crystallization at 200 °C; thus, sufficient E-MA-graf t-PLA copolymers are in situ formed in the blend melts. The E-MAgraf t-PLA can substantially improve the melt stability of SCPLA matrix as efficient compatibilizer capable of preventing the complete decoupling of PLLA and PDLA chains from their helical chain pairs by promoting the interchain interactions, which makes it possible to fabricate highly crystalline products with exclusive SC crystallites through injection molding of the PLLA/PDLA/E-MA-GMA blends. On the other hand, some E-MA-graf t-PLA chains localized at the blend interface can also greatly enhance the interfacial adhesion as interfacial enhancer. Consequently, injection-

■

REFERENCES

(1) Miller, S. A. Sustainable Polymers: Replacing Polymers Derived from Fossil Fuels. Polym. Chem. 2014, 5 (9), 3117−3118. (2) Schneiderman, D. K.; Hillmyer, M. A. 50th Anniversary Perspective: There Is a Great Future in Sustainable Polymers. Macromolecules 2017, 50 (10), 3733−3750. (3) Lim, L. T.; Auras, R.; Rubino, M. Processing Technologies for Poly(lactic acid). Prog. Polym. Sci. 2008, 33 (8), 820−852. (4) Reddy, M. M.; Vivekanandhan, S.; Misra, M.; Bhatia, S. K.; Mohanty, A. K. Biobased Plastics and Bionanocomposites: Current Status and Future Ppportunities. Prog. Polym. Sci. 2013, 38 (10−11), 1653−1689. (5) Farah, S.; Anderson, D. G.; Langer, R. Physical and Mechanical Properties of PLA, and Their Functions in Widespread Applications A Comprehensive Review. Adv. Drug Delivery Rev. 2016, 107, 367− 392. (6) Nagarajan, V.; Zhang, K.; Misra, M.; Mohanty, A. K. Overcoming the Fundamental Challenges in Improving the Impact Strength and Crystallinity of PLA Biocomposites: Influence of Nucleating Agent and Mold Temperature. ACS Appl. Mater. Interfaces 2015, 7 (21), 11203−11214. (7) Nagarajan, V.; Mohanty, A. K.; Misra, M. Perspective on Polylactic Acid (PLA) based Sustainable Materials for Durable Applications: Focus on Toughness and Heat Resistance. ACS Sustainable Chem. Eng. 2016, 4 (6), 2899−2916. K


Article

Macromolecules

High-Molecular-Weight Stereo Diblock Poly(lactic acid)s. J. Phys. Chem. B 2015, 119 (44), 14270−14279. (27) Jing, Z. X.; Shi, X. T.; Zhang, G. C.; Li, J.; Li, J. W.; Zhou, L. S.; Zhang, H. M. Formation, Structure and Promoting Crystallization Capacity of Stereocomplex Crystallite Network in the Poly(lactide) Blends based on Linear PLLA and PDLA with Different Structures. Polymer 2016, 92, 210−221. (28) Sakamoto, Y.; Tsuji, H. Stereocomplex Crystallization Behavior and Physical Properties of Linear 1-Arm, 2-Arm, and Branched 4-Arm Poly(L-lactide)/Poly(D-lactide) Blends: Effects of Chain Directional Change and Branching. Macromol. Chem. Phys. 2013, 214 (7), 776−786. (29) Bao, J. N.; Han, L. L.; Shan, G. R.; Bao, Y. Z.; Pan, P. J. Preferential Stereocomplex Crystallization in Enantiomeric Blends of Cellulose Acetate-g-Poly(lactic acid)s with Comb like Topology. J. Phys. Chem. B 2015, 119 (39), 12689−12698. (30) Ma, P. M.; Jiang, L.; Xu, P. W.; Dong, W. F.; Chen, M. Q.; Lemstra, P. J. Rapid Stereocomplexation between Enantiomeric Comb-Shaped Cellulose-g-poly(L-lactide) Nanohybrids and Poly(Dlactide) from the Melt. Biomacromolecules 2015, 16 (11), 3723− 3729. (31) Isono, T.; Kondo, Y.; Ozawa, S.; Chen, Y. G.; Sakai, R.; Sato, S.; Tajima, K.; Kakuchi, T.; Satoh, T. Stereoblock-like Brush Copolymers Consisting of Poly(L-lactide) and Poly(D-lactide) Side Chains along Poly(norbornene) Backbone: Synthesis, Stereocomplex Formation, and Structure-Property Relationship. Macromolecules 2014, 47 (20), 7118−7128. (32) Brzezinski, M.; Boguslawska, M.; Ilcikova, M.; Mosnacek, J.; Biela, T. Unusual Thermal Properties of Polylactides and Polylactide Stereocomplexes Containing Polylactide-Functionalized Multi-Walled Carbon Nanotubes. Macromolecules 2012, 45 (21), 8714−8721. (33) Han, L. L.; Pan, P. J.; Shan, G. R.; Bao, Y. Z. Stereocomplex Crystallization of High-Molecular-Weight Poly(L-lactic acid)/Poly(D-lactic acid) Racemic Blends Promoted by a Selective Nucleator. Polymer 2015, 63, 144−153. (34) Xiong, Z. J.; Zhang, X. Q.; Wang, R.; de Vos, S.; Wang, R. Y.; Joziasse, C. A. P.; Wang, D. J. Favorable Formation of Stereocomplex Crystals in Poly(L-lactide)/Poly(D-lactide) Blends by Selective Nucleation. Polymer 2015, 76, 98−104. (35) Bao, R. Y.; Yang, W.; Wei, X. F.; Xie, B. H.; Yang, M. B. Enhanced Formation of Stereocomplex Crystallites of High Molecular Weight Poly(L-lactide)/Poly(D-lactide) Blends from Melt by Using Poly(ethylene glycol). ACS Sustainable Chem. Eng. 2014, 2 (10), 2301−2309. (36) Zhu, J.; Na, B.; Lv, R. H.; Li, C. Enhanced Stereocomplex Formation of High-Molecular-Weight Polylactides by Gelation in an Ionic Liquid. Polym. Int. 2014, 63 (6), 1101−1104. (37) Pan, P. J.; Bao, J. N.; Han, L. L.; Xie, Q.; Shan, G. R.; Bao, Y. Z. Stereocomplexation of High-Molecular-Weight Enantiomeric Poly(lactic acid)s Enhanced by Miscible Polymer Blending with Hydrogen Bond Interactions. Polymer 2016, 98, 80−87. (38) Bao, J. N.; Xue, X. J.; Li, K.; Chang, X. H.; Xie, Q.; Yu, C. T.; Pan, P. J. Competing Stereocomplexation and Homocrystallization of Poly(L-lactic acid)/Poly(D-lactic acid) Racemic Mixture: Effects of Miscible Blending with Other Polymers. J. Phys. Chem. B 2017, 121 (28), 6934−6943. (39) Fukushima, K.; Chang, Y. H.; Kimura, Y. Enhanced Stereocomplex Formation of Poly(L-lactic acid) and Poly(D-lactic acid) in the Presence of Stereoblock Poly(lactic acid). Macromol. Biosci. 2007, 7 (6), 829−835. (40) Jing, F.; Hillmyer, M. A. A Bifunctional Monomer Derived from Lactide for Toughening Polylactide. J. Am. Chem. Soc. 2008, 130 (42), 13826−13827. (41) Liu, H. L.; Zhang, T. T.; Cai, Y.; Deng, S. H.; Bai, D. Y.; Bai, H. W.; Zhang, Q.; Fu, Q. Towards Polylactide/Core-Shell Rubber Blends with Balanced Stiffness and Toughness via the Formation of Rubber Particle Network with the Aid of Stereocomplex Crystallites. Polymer 2018, 159, 23−31.

(8) Tsuji, H. Poly(lactide) Stereocomplexes: Formation, Structure, Properties, Degradation, and Applications. Macromol. Biosci. 2005, 5 (7), 569−597. (9) Longo, J. M.; DiCiccio, A. M.; Coates, G. W. Poly(propylene succinate): A New Polymer Stereocomplex. J. Am. Chem. Soc. 2014, 136, 15897−15900. (10) Tsuji, H. Poly(lactic acid) Stereocomplexes: A Decade of Progress. Adv. Drug Delivery Rev. 2016, 107, 97−135. (11) Ikada, Y.; Jamshidi, K.; Tsuji, H.; Hyon, S. H. Stereocomplex Formation between Enantiomeric Poly(Lactides). Macromolecules 1987, 20 (4), 904−906. (12) Tsuji, H.; Ikada, Y. Stereocomplex Formation between Enantiomeric Poly(lactic acid)s. XI. Mechanical Properties and Morphology of Solution-Cast Films. Polymer 1999, 40 (24), 6699− 6708. (13) Zhang, X. W.; Nakagawa, R.; Chan, K. H. K.; Kotaki, M. Mechanical Property Enhancement of Polylactide Nanofibers through Optimization of Molecular Weight, Electrospinning Conditions, and Stereocomplexation. Macromolecules 2012, 45, 5494−5500. (14) Zhang, Z. C.; Sang, Z. H.; Huang, Y. F.; Ru, J. F.; Zhong, G. J.; Ji, X.; Wang, R. Y.; Li, Z. M. Enhanced Heat Deflection Resistance via Shear Flow-Induced Stereocomplex Crystallization of Polylactide Systems. ACS Sustainable Chem. Eng. 2017, 5 (2), 1692−1703. (15) Tsuji, H. In Vitro Hydrolysis of Blends from Enantiomeric Poly(lactide)s Part 1. Well-Stereo-Complexed Blend and NonBlended Films. Polymer 2000, 41 (10), 3621−3630. (16) Andersson, S. R.; Hakkarainen, M.; Inkinen, S.; Sodergard, A.; Albertsson, A. C. Polylactide Stereocomplexation Leads to Higher Hydrolytic Stability but More Acidic Hydrolysis Product Pattern. Biomacromolecules 2010, 11 (4), 1067−1073. (17) Bai, H. W.; Deng, S. H.; Bai, D. Y.; Zhang, Q.; Fu, Q. Recent Advances in Processing of Stereocomplex-Type Polylactide. Macromol. Rapid Commun. 2017, 38 (23), 1700454. (18) Samuel, C.; Cayuela, J.; Barakat, I.; Muller, A. J.; Raquez, J. M.; Dubois, P. Stereocomplexation of Polylactide Enhanced by Poly(methyl methacrylate): Improved Processability and Thermomechanical Properties of Stereocomplexable Polylactide-Based Materials. ACS Appl. Mater. Interfaces 2013, 5 (22), 11797−11807. (19) Tsuji, H.; Ikada, Y. Stereocomplex Formation between Enantiomeric Poly(Lactic Acid)S 0.9. Stereocomplexation from the Melt. Macromolecules 1993, 26 (25), 6918−6926. (20) He, Y.; Xu, Y.; Wei, J.; Fan, Z. Y.; Li, S. M. Unique Crystallization Behavior of Poly(L-lactide)/Poly(D-lactide) Stereocomplex Depending on Initial Melt States. Polymer 2008, 49 (26), 5670−5675. (21) Pan, P. J.; Han, L. L.; Bao, J. N.; Xie, Q.; Shan, G. R.; Bao, Y. Z. Competitive Stereocomplexation, Homocrystallization, and Polymorphic Crystalline Transition in Poly(L-lactic acid)/Poly(Dlactic acid) Racemic Blends: Molecular Weight Effects. J. Phys. Chem. B 2015, 119 (21), 6462−6470. (22) Bai, H. W.; Liu, H. L.; Bai, D. Y.; Zhang, Q.; Wang, K.; Deng, H.; Chen, F.; Fu, Q. Enhancing the Melt Stability of Polylactide Stereocomplexes using a Solid-State Cross-Linking Strategy during a Melt-Blending Process. Polym. Chem. 2014, 5 (20), 5985−5993. (23) Huang, Y. F.; Zhang, Z. C.; Li, Y.; Xu, J. Z.; Xu, L.; Yan, Z.; Zhong, G. J.; Li, Z. M. The Role of Melt Memory and Template Effect in Complete Stereocomplex Crystallization and Phase Morphology of Polylactides. Cryst. Growth Des. 2018, 18 (3), 1613−1621. (24) Hirata, M.; Kimura, Y. Thermo Mechanical Properties of Stereoblock Poly(lactic acid)s with Different PLLA/PDLA Block Compositions. Polymer 2008, 49 (11), 2656−2661. (25) Biela, T.; Duda, A.; Penczek, S. Enhanced Melt Stability of Star-Shaped Stereocomplexes as Compared with Linear Stereocomplexes. Macromolecules 2006, 39 (11), 3710−3713. (26) Han, L. L.; Shan, G. R.; Bao, Y. Z.; Pan, P. J. Exclusive Stereocomplex Crystallization of Linear and Multiarm Star-Shaped L


Article

Macromolecules (42) Zhao, T. H.; Yuan, W. Q.; Li, Y. D.; Weng, Y. X.; Zeng, J. B. Relating Chemical Structure to Toughness via Morphology Control in Fully Sustainable Sebacic Acid Cured Epoxidized Soybean Oil Toughened Polylactide Blends. Macromolecules 2018, 51 (5), 2027− 2037. (43) Ge, T.; Grest, G. S.; Robbins, M. O. Structure and Strength at Immiscible Polymer Interfaces. ACS Macro Lett. 2013, 2 (10), 882− 886. (44) Liu, H. L.; Bai, H. W.; Bai, D. Y.; Liu, Z. W.; Zhang, Q.; Fu, Q. Design of High-Performance Poly(l-lactide)/Elastomer Blends through Anchoring Carbon Nanotubes at the Interface with the Aid of Stereocomplex Crystallization. Polymer 2017, 108, 38−49. (45) Oyama, H. I. Super-Tough Poly(lactic acid) Materials: Reactive Blending with Ethylene Copolymer. Polymer 2009, 50 (3), 747−751. (46) Liu, H. Z.; Chen, F.; Liu, B.; Estep, G.; Zhang, J. W. Super Toughened Poly(lactic acid) Ternary Blends by Simultaneous Dynamic Vulcanization and Interfacial Compatibilization. Macromolecules 2010, 43 (14), 6058−6066. (47) Liu, H. Z.; Song, W. J.; Chen, F.; Guo, L.; Zhang, J. W. Interaction of Microstructure and Interfacial Adhesion on Impact Performance of Polylactide (PLA) Ternary Blends. Macromolecules 2011, 44 (6), 1513−1522. (48) Wu, M.; Wu, Z. Q.; Wang, K.; Zhang, Q.; Fu, Q. Simultaneous the Thermodynamics Favorable Compatibility and Morphology to Achieve Excellent Comprehensive Mechanics in PLA/OBC Blend. Polymer 2014, 55 (24), 6409−6417. (49) Mauck, S. C.; Wang, S.; Ding, W. Y.; Rohde, B. J.; Fortune, C. K.; Yang, G. Z.; Ahn, S. K.; Robertson, M. L. Biorenewable Tough Blends of Polylactide and Acrylated Epoxidized Soybean Oil Compatibilized by a Polylactide Star Polymer. Macromolecules 2016, 49 (5), 1605−1615. (50) Mazidi, M. M.; Edalat, A.; Berahman, R.; Hosseini, F. S. Highly-Toughened Polylactide- (PLA-) Based Ternary Blends with Significantly Enhanced Glass Transition and Melt Strength: Tailoring the Interfacial Interactions, Phase Morphology, and Performance. Macromolecules 2018, 51 (11), 4298−4314. (51) Zolali, A. M.; Favis, B. D. Toughening of Cocontinuous Polylactide/Polyethylene Blends via an Interfacially Percolated Intermediate Phase. Macromolecules 2018, 51 (10), 3572−3581. (52) Chen, D. P.; Wang, H. T.; Li, Y. Reactive Compatibilization: Formation of Double-Grafted Copolymers by In Situ Binary Grafting and Their Compatibilization Effect. ACS Appl. Mater. Interfaces 2017, 9 (38), 33091−33099. (53) Bai, D. Y.; Liu, H. L.; Bai, H. W.; Zhang, Q.; Fu, Q. LowTemperature Sintering of Stereocomplex-Type Polylactide Nascent Powder: Effect of Crystallinity. Macromolecules 2017, 50 (19), 7611− 7619. (54) Bai, D. Y.; Diao, X. Y.; Ju, Y. L.; Liu, H. L.; Bai, H. W.; Zhang, Q.; Fu, Q. Low-Temperature Sintering of Stereocomplex-Type Polylactide Nascent Powder: The Role of Optical Purity in Directing the Chain Interdiffusion and Cocrystallization across the Particle Interfaces. Polymer 2018, 150, 169−176. (55) Luo, Y. L.; Ju, Y. L.; Bai, H. W.; Liu, Z. W.; Zhang, Q.; Fu, Q. Tailor-Made Dispersion and Distribution of Stereocomplex Crystallites in Poly(L-lactide)/Elastomer Blends toward Largely Enhanced Crystallization Rate and Impact Toughness. J. Phys. Chem. B 2017, 121 (25), 6271−6279. (56) Dong, W. Y.; Wang, H. T.; Ren, F. L.; Zhang, J. Q.; He, M. F.; Wu, T.; Li, Y. J. Dramatic Improvement in Toughness of PLLA/ PVDF Blends: the Effect of Compatibilizer Architectures. ACS Sustainable Chem. Eng. 2016, 4 (8), 4480−4489. (57) Bao, R. Y.; Yang, W.; Jiang, W. R.; Liu, Z. Y.; Xie, B. H.; Yang, M. B.; Fu, Q. Stereocomplex Formation of High-Molecular-Weight Polylactide: A Low Temperature Approach. Polymer 2012, 53 (24), 5449−5454. (58) Song, Y.; Wang, D. J.; Jiang, N.; Gan, Z. H. Role of PEG Segment in Stereocomplex Crystallization for PLLA/PDLA-b-PEG-bPDLA Blends. ACS Sustainable Chem. Eng. 2015, 3 (7), 1492−1500.

(59) Yang, C. F.; Huang, Y. F.; Ruan, J.; Su, A. C. Extensive Development of Precursory Helical Pairs Prior to Formation of Stereocomplex Crystals in Racemic Polylactide Melt Mixture. Macromolecules 2012, 45 (2), 872−878. (60) Tashiro, K.; Kouno, N.; Wang, H.; Tsuji, H. Crystal Structure of Poly(lactic acid) Stereocomplex: Random Packing Model of PDLA and PLLA Chains As Studied by X-ray Diffraction Analysis. Macromolecules 2017, 50 (20), 8048−8065. (61) Wu, S. H.; Wang, F. Y.; Ma, C. C. M.; Chang, W. C.; Kuo, C. T.; Kuan, H. C.; Chen, W. J. Mechanical, Thermal and Morphological Properties of Glass Fiber and Carbon Fiber Reinforced Polyamide-6 and Polyamide-6/Clay Nanocomposites. Mater. Lett. 2001, 49 (6), 327−333. (62) Zang, C. G.; Zhu, X. D.; Jiao, Q. J. Enhanced Mechanical and Electrical Properties of Nylon-6 Composite by Using Carbon Fiber/ Graphene Multiscale Structure as Additive. J. Appl. Polym. Sci. 2015. (63) Kausar, A. A. Review of Filled and Pristine Polycarbonate Blends and Their Applications. J. Plast. Film Sheeting 2018, 34 (1), 60−97. (64) Deshmukh, G. S.; Peshwe, D. R.; Pathak, S. U.; Ekhe, J. D. A Study on Effect of Mineral Additions on the Mechanical, Thermal, and Structural Properties of Poly(butylene terephthalate) (PBT) Composites. J. Polym. Res. 2011, 18 (5), 1081−1090. (65) Arslan, C.; Dogan, M. The Effects of Silane Coupling Agents on the Mechanical Properties of Basalt Fiber Reinforced Poly(butylene terephthalate) Composites. Composites, Part B 2018, 146, 145−154.

M


Toward Supertough and Heat-Resistant Stereocomplex-Type

Recommend Documents