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Dramatic Improvement in Toughness of PLLA/PVDF Blends: the Effect of Compatibilizer Architectures Wenyong Dong, Hengti Wang, Fanglu Ren, Junqing Zhang, Meifeng He, Tao Wu, and Yongjin Li ACS Sustainable Chem. Eng., Just Accepted Manuscript • DOI: 10.1021/ acssuschemeng.6b01420 • Publication Date (Web): 27 Jun 2016 Downloaded from http://pubs.acs.org on July 1, 2016

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ACS Sustainable Chemistry & Engineering

Dramatic Improvement in Toughness of PLLA/PVDF Blends: the Effect of Compatibilizer Architectures

Wenyong Dong, Hengti Wang, Fanglu Ren, Junqing Zhang, Meifeng He, Tao Wu, Yongjin Li*

College of Material, Chemistry and Chemical Engineering, Hangzhou Normal University, No.16 Xuelin Road, Hangzhou 310036, P.R. China

Corresponding author. Email: [email protected] TEL: 86-571-2886-7206; FAX: 86-571-2886-7899

ACS Paragon Plus Environment


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Abstract Reactive comb (RC) polymers with different molecular architectures have been successfully synthesized by copolymerizing methyl methacrylate (MMA), glycidyl methacrylate (GMA) and a series of MMA macromer with different molecular weights. The prepared RC polymers with different lengths of side chains were applied as compatibilizers in an immiscible poly(L-lactic acid)/poly(vinylidene fluoride) (PLLA/PVDF) blend and it was found that the RC polymers with moderate length of side chain (e.g. Mn = 4800 g mol-1) displayed better compatibilizing efficiencies than RC polymers with short side chain (e.g. Mn = 2400 g mol-1) and let alone the Reactive Linear (RL) polymers without side chains. The thus obtained PLLA/PVDF blends with PLLA as matrix will provide excellent adhesion with the surfaces of the metals and cells and it was found that RC polymers endowed the PLLA/PVDF blend with excellent toughness and the RC-compatibilized blend could be uniaxially stretched to a maximum draw ratio of 5 at room temperature (about 22 oC). FT-IR and XRD results showed that the nonpolar α phase of PVDF was completely transformed into the piezoelectric β phase (more than 95%) during the stretching.

Keywords: Reactive compatibilization, Piezoelectric, Poly(lactic acid), Poly(vinylidene fluoride), Interface


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Introduction Reactive compatibilization by the in situ formation of block or graft copolymers at the interface of the immiscible polymer blends is a cost effective way to fabricate high performance materials with the combination of the properties of the constituents; especially, reactive formed graft copolymers by the end-functionalized chains in one phase with a multifunctional chains in the other phase is more appealing in industrial application and numerous investigations have been devoted to this research area.[1,2] Poly(L-lactic acid) (PLLA) is a semicrystalline thermoplastic polymer and has attracted much attention during the past few years, because it is derived from sustainable sources (e.g. corn starch, tapioca roots) and is often considered as a green alternative to petroleum-based polymers.[3,4] Although the mechanical properties of PLLA was similar to that of polystyrene, the inherent brittleness inevitably limited the applications of PLLA in many areas. To improve the toughness of PLLA, both renewable (e.g. poly(butylene adipate-co-terephthalate) (PBAT)[5] and polycaprolactone (PCL)[6]) and petroleum-based polymers (e.g. polycarbonate (PC)[7] and polyethylene (PE)[8]) were reported to be blended with PLLA and the latter was thought to be a more cost effective way. PLLA was well-known for its biocompatible and biodegradable properties and less well-known to be a piezoelectric polymer.[9-11] It is because PLLA has asymmetric carbon atoms, which displays a helical orientation that it exhibits shear piezoelectricity, while its piezoelectric constant is only one-third of that of poly(vinylidene fluoride) (PVDF). PVDF is also a semicrystalline thermoplastic polymer. It has many applications from the biological and biomedical fields to the electrical and electronic fields and this is due to its electroactive behavior (e.g. high piezo- and pyro-electric), good biocompatibility, heat and chemical resistance and excellent mechanical properties, etc. PVDF is also a kind of polycrystalline polymers and amongst the five different crystalline phases (α, β, γ, δ and ε), the β phase has the best piezoelectric and pyroelectric properties, because its molecular structure displays an all-trans planar zigzag (TTTT) conformation.[12,13] Many methods have been adopted to obtain the β phase PVDF, such as uniaxial or biaxial stretching,[14,15] addition of nanofillers (e.g. carbon nanotube,[16] graphene oxide,[17] halloysite nanotube[18]), or room temperature ionic liquid (RTIL) decorated carbon nanotube,[19] blending with other polymers (e.g. PMMA,[20] Polyaniline



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(PANI)[21]) and crystallization from polar solvents.[22] Although stretching induced `transformation of PVDF from α to β was a convenient method, such a process was often accomplished above at least 60 oC, because of the limited elongation of PVDF at room temperature. We noted that the group of Maiti recently reported a series of works on improving the piezoelectricity of PVDF;[23,24] one of them was by adding Bis(hydroxyethyl)methyltallow ammonium ion exchanged montmorillonite and the obtained nanocomposites by the solution processed method displayed a maximum elongation at break of about 450% at room temperature and nearly 1000% at 90 oC, respectively, during which process, the α phase was completely transformed into the β phase.[23] Accelerating the development of the renewable energy (e.g. solar, wind and energy produced by the human body) acquiring from the environment is an urgent need for the sustainable future of human race. Piezoelectric materials are perfect candidates in the energy harvesting system for they can convert mechanical energy into electrical energy, i.e., the biomechanical energy from human motion is potentially transformed into electric power.[25,26] Although PVDF has technological importance, the inherent low surface energy prohibits its adhesion with other materials like metals and cells. In order to improve the adhesion of PVDF, some methods like chemical decoration and plasma treatment were applied and these inevitably resulted in the destruction of the molecular chains of PVDF.[27-30] We noted that the group of Loos recently designed and synthesized a series of

poly(L-lactide)-b-poly(vinylidene

fluoride)-b-poly(L-lactide)

(PLLA-b-PVDF-b-PLLA)

triblock copolymers and systematically investigatedthe relationships between the crystallization behaviors of the components and the block ratios.[31] This kind of block copolymers skillfully combined the advantages of both PLLA and PVDF and was a promising electroactive material in energy harvesting devices. Almost at the same time, we provided a more convenient way to prepare PLLA/PVDF materials with excellent toughness by reactive compatibilization. A kind of reactive comb (RC) polymers was applied as compatibilizers in the PLLA/PVDF blend, in which even if the weight ratio of PLLA/PVDF was 50/50, PLLA could exist as the matrix and PVDF as the dispersed phase, because PVDF had higher viscosity and lower volume fraction than that of PLLA. The polar PLLA matrix could provide excellent adhesion of the blend with other materials and we found that the toughness of the blend was significantly improved after compatibilization and the maximum fraction strain was nearly 300%.[32] The RC polymers was synthesized by copolymerizing methyl


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methacrylate (MMA), glycidyl methacrylate (GMA) and MMA macromer.[32,33] Herein, the weight ratio of these three monomers was further optimized and at the same time, the Mn of the MMA monomer was varied; thus we obtained a series of RC polymers with different lengths of side chains. Then we investigated the relationships between the molecular structure of the PMMA-type compatibilizers and the toughness of the compatibilized PLLA/PVDF blends. We further evaluated the β phase content of PVDF in the RL- or RC-compatibilized PLLA/PVDF blends after uniaxial drawing at room temperature. Experimental section Materials PVDF (KF850) and PLLA (3001D) homopolymers were purchased from Kureha Chemicals and Nature works, respectively. The RC polymers were synthesized by copolymerizing MMA and GMA with three kinds of MMA macromer with different molecular weights. MMA macromer was synthesized by two steps. Firstly, carboxyl-terminated PMMA (PMMA-COOH) was synthesized by the telomerization of MMA with 4, 4’-azobis(4-cyanovaleric acid)(ACVA) as initiator and thioglycolic acid (TAC) as chain transfer agent; the molecular weight of PMMA-COOH could be tuned by the molar ratio ACVA/TAC. As shown in Figure S2, it was found that the molecular weight of PMMA-COOH increased with the decreasing of the molar ratio of ACVA/TAC and thus we synthesized three typical PMMA-COOH with different molecular weight (different length of molecular chains). Secondly, PMMA-COOH reacted with GMA and the carboxyl groups were changed into double bonds, the molecular characterizations of the MMA macromers were shown in Table S1. In our previous report, the weight ratio of MMA macromer/GMA/MMA was 1/1/8, while in the present investigation, the weight ratio was changed to 1/2/7. The 1H NMR spectrum of RC polymers was shown in Figure S1 and compared to our former report, the increase of the intensity of the peaks corresponding to the epoxy groups was ascribed to the increase of the weight percent of GMA in RC polymers.[32,33] As shown in Table 1, three kinds of RC polymers with increasing lengths of side chains were synthesized. A linear polymer by copolymerizing GMA and MMA at a weight of 2/8 and a kind of RC polymer (MMA macromer/GMA/MMA = 1/1/8) reported previously were also synthesized for comparison. The molecular characterizations of the RC and RL polymers in Table 1 showed that the Mn of both the RL and RC polymers were all the same, in other words, the backbone had the same length.



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Sample ID

Table 1. Molecular characterization of RL and RC polymers Composition (%) Mn

Mw/Mn (g mol-1) Macromer GMA MMA L-0-1-9 a 0 10 90 22000 2.0 L-0-2-8 0 20 80 21000 1.9 b C-1-2-7(S-2400) 10 20 70 20000 2.1 C-1-2-7(S-4800) 10 20 70 20000 2.0 C-1-2-7(S-6300) 10 20 70 19000 2.0 C-1-1-8(S-2400) 10 10 80 21000 1.9 a L is short for Linear and the numbers connected by dash is the weight ratio of MMA macromer/GMA/MMA b C is short for Comb, S in bracket is short for side chain and the number in bracket is the number average molecular weight of the MMA macromer PLLA/PVDF blend preparation PLLA, PVDF and our lab-made PMMA-type compatibilizers were dried in a vacuum oven at 80 oC overnight before use. The compatibilized blends were prepared in a batch mixer (Haake Polylab QC) equipped with two counter-rotating roller blades at a speed of 50 rpm for 10 min at 190 oC. The weight ratio of PLLA/PVDF was fixed to 50/50, and the weight content of the RL or RC in the blends was 1, 3 and 5 wt % with respect to the total weight amount of the PLLA/PVDF blends. After blending, the samples were compression-molded under 10 MPa at 200 oC into sheets with thickness of 0.5 mm and the obtained sheets were used directly in the subsequent characterizations. Characterizations 1

H Nuclear magnetic resonance (1H NMR) spectra were recorded with a Bruker Avance 500

spectrometer at room temperature, operating at a frequency of 500 MHz, with deuterated chloroform as the solvent. Gel permeation chromatography (GPC) was carried out to determine the number-average molecular weight (Mn) and dispersity (Đ) using two MZ-Gel SD plus 10.0 µm bead-size columns (10E5 and 10E3 Å) and Optilab T-rEX detector. THF was used as the mobile phase at a flow rate of 1 mL min-1 at 35 °C. The system was calibrated with narrow molecular weight distribution polystyrene standards from 2000 to 106 g mol-1. The elution diagrams were analyzed using the ASTRA 6 software from Wyatt Technology. Fourier transform infrared (FT-IR) measurements by the attenuated total reflectance (ATR) method were performed on a Thermo Scientific Nicolet iS5 with iD5 ATR accessory. The spectra


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were scanned for 16 times at a resolution of 2 cm-1.The samples for FT-IR investigation were prepared by uniaxially drawing the compression-molded sheets at a crosshead speed of 10 mm min-1 to a draw ratio just before the sample break. X-ray diffraction (XRD) measurements were carried out on a Bruker D8 advance powder diffractometer, using a nickel filtered Cu Kα radiation (λ= 0.1542 nm). Data were collected at a scanning speed of 4° min-1 with a step interval of 0.04°. Differential scanning calorimetry (DSC) measurements were carried out from DSC Q2000 (TA Instrument) in an inert atmosphere of nitrogen at a heating or cooling rate of 10 °C min-1 with a temperature range from -50 to 200 °C. Scanning electron microscopy (SEM) images were recorded using a field emission scanning electron microscope (Hitachi S-4800) at an accelerating voltage of 5.0 KV. All the samples were fractured by immersing in liquid nitrogen for about 10 min. The fracture surface was then sputtered with a thin layer of gold. Particle size was countered for more than 300 particles. Atomic force microscopy (AFM) images were recorded using a SEIKO SPA-300 instrument in tapping mode. A silicon tip with a spring constant of 1.4 N m−1 purchased from SEIKO was used. The samples for AFM investigation were ultramicrotomed to obtain a flat surface. Transmission electron microscopy (TEM) was performed using a Hitachi HT-7700 instrument operating at an accelerating voltage of 80 kV. The blend samples were firstly ultramicrotomed to obtain a section with a thickness of about 70~100 nm and then stained with Ruthium tetroxide (RuO4) for 2 h. Tensile tests were carried out using an Instron universal material testing system (model 5966) at 22 °C at a crosshead speed of 10 mm min-1. The specimens were punched out of the compression-molded films into a dumbbell shape (18 mm gauge length, 0.5mm gauge thickness and 3 mm gauge width). Rheological measurements were performed on a Physica rheometer (MCR301, Anton Paar Instrument) at 200 °C in a nitrogen atmosphere. A parallel plate configuration was used with a plate diameter of 25 mm and a gap size of 1 mm. Frequency sweep was carried out from 500 to 0.01rad s-1 at a strain of 0.05 to ensure the sweep was in the linear viscoelastic region.



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Results PLLA/PVDF blends compatibilized by RL polymers: the effect of the GMA content In a previous report, we synthesized and compared the compatibilization efficiencies of RL and RC polymers; the RC polymers were synthesized by copolymerizing MMA macromer /GMA/MMA at a weight ratio of 1/1/8 and the weight ratio of GMA /MMA in RL polymers was 1/9.[32] In the present investigation, we firstly increased the weight ratio of GMA in RL polymers from 10 wt% to 20wt% to investigate the effect of GMA content on the morphology development of the PLLA/PVDF blend.[34,35] We found that the size of the dispersed phase decreased and the homogeneity of the dispersed phase increased, with the addition of the compatibilizers (Figure S3). As shown in Figure 1, the emulsification curve, which displayed the evolution of the number average diameter of the dispersed phase (dn) as a function of the concentration of the compatibilizer, indicated that although both curves had the same trend, the curve of L-0-2-8 descended more rapidly than that of L-0-1-9. During the reactive blending, the epoxide groups of RL polymers reacted with the carboxyl groups of PLLA at the interface of the immiscible blend to reduce the interfacial tension and to inhibit the coalescence of the dispersed phase. With the increasing of the GMA content in RL polymers, the probability of interfacial reaction between the epoxy and carboxyl groups was increased. So with the addition of only 1 wt% of L-0-2-8, the size of the dispersed phase decreased significantly and the emulsification curves leveled off when the concentration of the RL polymers was higher than 3 wt%, which meant the saturation effect of the compatibilizers in-situ formed at the interface.[36]

Figure 1. The number average diameter (dn) of the dispersed phase in PLLA/PVDF (50/50) blends versus weight % of L-0-1-9 (A) and L-0-2-8 (B)


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PLLA/PVDF blends compatibilized by RC polymers: the effect of the length of the side chains

Figure 2. AFM images of PLLA/PVDF (50/50) blends uncompatibilized (A) and compatibilized by 3wt% of (B) L-0-2-8; (C) C-1-2-7(S-2400); (D, E) C-1-2-7(S-4800); (F) C-1-2-7(S-6300) In the present investigation, the content of GMA in RC polymers was further increased from 10 wt% to 20 wt% and a series of RC polymers with the same length of the backbone and different lengths of side chains were synthesized. As shown in Figure 2, the dark phase in the micrographs corresponded to PLLA and the brighter phase to PVDF. Compared to RL polymer (L-0-2-8) in Figure 2 B, the size of the dispersed phase compatibilized by RC polymers was significantly reduced (Figure 2 C-F). As the length of the side chains increased from 2400 g mol-1 to 6300 g mol-1, the morphology evolved from a little irregular big particles, ellipsoidal fibers and finally to homogeneous small particles. The TEM results (Figure S3) were in consistent with those of AFM on the morphology evolution with the variation of the compatibilizer’s topology. Macosko et al once proposed a lacing/sheeting mechanism for the morphology development of



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polymer blend and the morphology evolution in our system could be tentatively explained as follows.[37-39] The PLLA phase melted first, and then the high viscous PVDF phase was gradually stretched into sheets by the less viscous PLLA phase as temperature increased. Under shear condition, holes began to form in the PVDF sheets and the sheets were torn up and stretched into fibers, which finally broke up into small particles. Thus the irregular big particles, ellipsoidal fibers and the homogeneous small particles mentioned above in our system might be related with this process. Another explanation for the morphology evolution could be by the capillary number (Ca), which was important for the deformation and break-up of drops in shear field. Ca was defined as the ratio of the shear force to the interfacial force:

Ca =

ηγ& R σ

(1)

in which η was the matrix viscosity, γ& the shear rate and σ the equilibrium interfacial tension. The in-situ formed compatibilizers at the interface could reduce the interfacial tension and thus facilitate the drop deformation and breakup.[40,41] From the morphology discrepancies of the blends compatibilized by a series of RL or RC polymers (Figure 2B-F), we could estimate that the blend compatibilized by C-1-2-7 (S-4800) had the lowest interfacial tension, because the drops exhibited the largest deformation. The typical details of the interface were shown in Figure S5. Compared to the uncompabitilized blend which displayed a sharp interface, the compatibilized blends displayed an intermediate region of medium darkness, which was derived from the in-situ formed compatibilizers at the interface. Recently, Zhang et al found that the entanglement molecular weight between the dissimilar chains in the PMMA/PVDF blend (Me12) was less than that of the blend component; it meant that the PMMA and PVDF chains were more likely to interact with each other due to the specific interaction between PMMA and PVDF. When the weight ratio of PVDF/PMMA was 90/10, the Me12 was 5700 g mol-1 and this value had a trend to decrease with the decreasing of the weight ratio of PMMA.[42] We could deduce that Me12 in the reactive compatibilized PLLA/PVDF system was slightly less than 5700 g mol-1, because the weight ratio of the PVDF/PMMA-type compatibilizer was 100/6. So we could understand why C-1-2-7(S-4800) exhibited the greatest ability to reduce the interfacial tension. Compared to C-1-2-7(S-2400), the molecular weight of the side chains in C-1-2-7(S-4800) was 4800 g mol-1,


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which was more close to the entanglement weight between PMMA and PVDF (Me12 = 5700 g mol-1 when the PVDF/PMMA = 100/10). Mechanical properties

Figure 3. Representative stress-strain curves of (a) PLLA/PVDF (50/50), (b) PLLA/PVDF/L-0-2-8 (50/50/3), (c) PVDF and (d) PLLA/PVDF/C-1-2-7(S-4800) (50/50/3) The stress-strain curves and the tensile properties of PLLA/PVDF blends uncompatibilized and compatibilized by RL or RC polymers were shown in Figure 3, Figure S6-8 and Table 2, respectively. The macroscopic pictures of some representative tensile bars were also provided in Figure S9 and it was found that the compatibilized blends all exhibited necking, while the uncompatibilized blend fractured without yielding. As summerized in Table 2, the fracture strain was 97.4% when the blend was compatibilized by 3 wt% of L-0-2-8. Compared with that by RL polymers, the fracture strain of the blends compatibilized by RC polymers was at least three times higher than that by RL. For example, when the length of the side chains increased from 2400 to 4800 g mol-1, the fracture strain of the blends compatibilized by 3wt% of C-1-2-7(S-4800) increased from 361% to 495%, almost five times higher than that by 3wt% of L-0-2-8 (Figure 3). It must be noticed that the wt% of GMA were the same in both C-1-2-7(S-4800) and L-0-2-8 and it was only the existence of the side chains that significantly improved the ductility of the blends. An interesting difference between the blends compatibilized by RL or RC polymers was that the stress-strain curves showed a strain hardening when the strain was over 220% for the RC-compatibilized blends. Especially, the strain hardening was pronounced for the blends compatibilized by 3wt% of C-1-2-7(S-4800) and the fracture stress was almost identical to the



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yield stress. Table 2. Mechanical properties of neat PLLA, neat PVDF and PLLA/PVDF (50/50) blends Fracture Fracture Modulus Yield stress Sample stress strain (GPa) (MPa) (MPa) (%) PLLA 1.67 61.2 51.8 9.4 PVDF 1.25 47.9 35.3 254 PLLA/PVDF 1.56 61.2 5.9 (50/50)a PLLA/PVDF/L-0-2-8 1.45 68.5 41.5 59 (50/50/1)b PLLA/PVDF/L-0-2-8 1.4 65.5 38.9 97.4 (50/50/3) PLLA/PVDF/L-0-2-8 1.54 67.1 40.8 127 (50/50/5) PLLA/PVDF/C-1-1-8(S-2400) 1.56 67.7 47.3 301 (50/50/3) PLLA/PVDF/C-1-2-7(S-2400) 1.31 54.8 43.3 362 (50/50/3) PLLA/PVDF/C-1-2-7(S-4800) 1.21 56.4 55.5 495 (50/50/3) PLLA/PVDF/C-1-2-7(S-6300) 1.26 56.2 45.8 393 (50/50/3) a The numbers in bracket was the weight ratio of PLLA/PVDF b The numbers in bracket was the weight ratio of PLLA/PVDF/compatibilizer Rheological properties

Figure 4. Complex viscosity (‫׀‬η*‫ )׀‬as a function of frequency at 200 °C for neat components, uncompatibilized and compatibilized blends by RL or RC polymers


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Figure 4 displayed the complex viscosity (η*) of both the compatibilized blends and the neat components as a function of frequencies. PLLA exhibited a Newtonian plateau from the low to the medium frequencies,[43] while PVDF displayed a shear thinning behavior, which might be ascribed to the polydispersity of its molecular weight.[44] The maximum shear rate ( γ& ) in the internal batch mixer was estimated to be about 50 s-1 and as shown in Figure 4, the viscosity ratio (a ratio of the dispersed phase viscosity to the matrix viscosity) was more than 1 at the shear rate of 50 s-1. Reactive compatibilization increased the viscosity of PLLA reaction due to the grafting reaction between the carboxyl end groups of PLLA and the epoxy groups of compatibilizers. Thus the viscosity ratio might decrease as the reaction time going on, which facilitated the PVDF dispersed phase to be stretched into fibers and finally into particles.[45-48] The complex viscosity of the uncompatbilied PLLA/PVDF blend was in between those of the neat components and the Newtonian plateau was observed only at low frequencies. This was consistent with the blend morphology that PLLA was existed as matrix, PVDF as dispersed phase and the interface between these two phases was weak. It should be noted that the complex viscosity of the blend compatibilized by RL polymers was higher than that of the uncompatibilized blends in the terminal region and was almost the same at high frequencies region, while the complex viscosity of the blend by RC polymers was significantly higher than that by RL polymers in the whole investigated range of frequencies. For example, the complex viscosity of RC compatibilized blend was about 3300 Pa s-1, almost two times higher than that by RL-compatibilized blend, which was about 1500 Pa s-1, at an angular frequency of 1 rad s-1. The storage modulus, loss modulus and loss tangent as a function of frequency were shown in Figure 5. The terminal behavior of PLLA followed the Maxwell model for linear and narrow-distribution polymers (G’∝ω2, G”∝ω); PVDF did not obey this behavior and this might be related to its polydispersed properties.[44] The G’ and G” of both the RL- and RC-compatibilized blends were much larger than the neat components and the uncompatibilized blend at low frequencies. Although both RL and RC-compatibilized blends showed a frequency-independent plateau in the terminal region, the G’ and G” of the RC-compatibilized blend were 5 and 2.4 times larger than that by RL polymers at the lowest frequency (0.01 rad s-1), respectively.[49] The low-frequency plateau of G’ displayed a non-terminal behavior, which was ascribed to the network formation of the PVDF dispersed phase and the



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differentiation of the G’ results implied that the elastic network-structure in RC-compatibilized blend was better developed and more dispersed phase particles were entrapped in this structure than that in RL-compatibilized blend. Compared to the RL-compatibilized blend, the tan δ peak of the RC-compatibilized blend shifted to higher frequencies with lower damping, which indicated that the elasticity was increased at low frequencies.[50-52]

Figure 5. Storage modulusG’ (a), loss modulusG” (b) and loss tangent Tan(δ) (c) as a function of


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frequency at 200oC for neat components, uncompatibilized and compatibilized blends by RL or RC polymers Thermal behaviors Both the PLLA and PVDF are semicrystalline polymers. Interestingly, in the PLLA and PVDF block copolymer system, the melting peaks of PLLA and PVDF in the DSC curves were separated from one the other,[31] while herein they were superimposed with each other because their very close melting points. It is therefore difficult to calculate the exact crystallinity of PVDF and PLLA by DSC (Figure S10). As shown in Figure 6 and 7, the melting peak temperatures (Tms) of both the PLLA and PVDF depressed gradually with the increasing of RL content, while the Tms were almost unaffected with the variation of the length of side chains in RC-compatibilized blends. In the RL-compatibilized blends, the cooling thermogram exhibited two peaks and the lower one corresponded to PLLA, the higher to PVDF, respectively. Compared to the uncompatibilized blend, the crystallization peak temperature (Tc) of PLLA in the compatibilized blend increased a little to higher temperature due to the crystalline PVDF as nucleating agent across PLLA/PVDF phase to promote the crystallization of PLLA. While with the increasing of RL content from 1 to 5 wt%, the Tcs of both the PLLA and PVDF phase in the compatibilized blends were gradually depressed. This indicated that the in-situ formed compatibilizers tend to interfere with the crystallization of both the PLLA and PVDF.[53-55] As shown in Figure 7 B, in the RC-compatibilized blends, the Tc of PLLA were nearly unaffected as a function of the PMMA side chain length, while the shape and intensity of the Tc corresponding to PVDF was significantly affected by the length of the side chains. The crystallization peak of PVDF was split into a main peak and a weak shoulder when the length of side chain was 2400 g mol-1. As the length increased to 4800 g mol-1, the intensity of the shoulder increased, while the intensity of the main peak decreased. With the further increase of the PMMA side chains to 6300 g mol-1, the intensity of the shoulder decreased again and the position of the main peak slightly shifted to higher temperature. It was conceivable that the PMMA side chains, which penetrated into the PVDF phase, significantly influenced the perfect crystallization of PVDF and some imperfect crystallites (i.e. needle-like) might be formed.[23]



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Figure 6. DSC thermograms of the first heating (A), cooling (B) and the second heating (C) of the PLLA/PVDF (50/50) blends compatibilized by different wt % of RL polymers

Figure 7. DSC thermograms of the first heating (A), cooling (B) and the second heating (C) of the PLLA/PVDF (50/50) blends compatibilized by 3 wt% of RC polymers with different lengths of side chains

α to β Phase Transformation in PVDF Blending PMMA with PVDF was one of the routes to obtain β-crystals of PVDF. The XRD results of the neat components, uncompatibilized and compatibilized blends by PMMA-type RL or RC polymers were shown in Figure S10. The unstretched samples (Figure S10 b-g) displayed the characteristic α-crystals of PVDF at 18.1o, 19.8o and 26.4o, corresponding to (020), (110) and (021) planes, respectively; 2θ values at 16.4o (Figure S10 a) was ascribed to (200)/(100) plane of PLLA. The FT-IR results (Figure S12) of the unstretched samples were in accordance with those of the XRD (Figure S11) and displayed the characteristic absorption bands at 765, 795 and 976 cm-1, respectively, corresponding to α-crystals of PVDF; [56] PLLA showed a weak and broad absorption at 754 cm-1, which was partly superimposed with the absorption of PVDF at 765 cm-1. The direct blending of PVDF and PLLA had been reported by us and some other authors.[57,58] In order to carry out the stretching at high temperature, PVDF was selected as the matrix and PLLA as the dispersed phase due to the poor thermal properties of PLLA, but this morphology


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with PVDF as the matrix had no benefits for the improvement of adhesion. It had been reported that temperature was a decisive factor during the drawing process of PVDF, because drawing at low temperature could not afford the molecular chains sufficient mobility to move, while at high temperature the already stretched chains tended to reorient to the α phase. For example, Xie et al systematically investigated the relationship between the drawing temperature and the β-phase content of PVDF in the PLLA/PVDF blend (the weight ratio of PLLA/PVDF was from 5/95 to 20/80); they found that the variations of the temperature significantly influenced the content of β-phase and when the drawing temperature increased from 60 oC to 100 oC, the content of β-phase decreased from about 90% to 60%.[57] The reactive compatibilization endowed the PLLA/PVDF blend with excellent toughness, even if the PLLA in the compatibilized blend existed as matrix, all of the compatibilized samples could be stretched at room temperature; the neat PVDF was also stretched at room temperature for comparison. It was found that the diffraction peaks corresponding to α phase were completely disappeared in RC-compatibilized samples; while in RL-compatibilized ones and neat PVDF, we could still observe the weak diffractions corresponding to α phase at 18.1o and 19.8o (Figure 8). FT-IR was applied to calculate the exact β-phase content in all the samples. As shown Figure 9, apart from the absorption bands of α phase, the bands at 840 cm-1 and 1279 cm-1 were characteristic of β-phase of PVDF. The β-phase content (F(β)) could be calculated by the following equation (2):

F (β ) =

Aβ 1.26 Aα + Aβ

(2)

in which Aα and Aβ were the areas of the absorption bands at 764 and 840 cm-1, characteristic of the α- and β-phases, respectively.[57] The calculated results were summarized in Table 3. It could be found that for neat PVDF and RL-compatibilized blend, only part of the α-phase were transformed into β-phase during stretching, while for RC-compatibilized blends, especially those RC polymers with long side chains (C-1-2-7(S-4800) and C-1-2-7(S-6300)), the transformations were almost complete. These results above indicated that high draw ratio of the compatibilized PLLA/PVDF blend at room temperature was readily attained by RC polymers, especially those with moderate length of side chains and the α phase of the PVDF was almost completely transformed into the β phase during the drawing process.



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Figure 8. XRD profiles for stretched (a) PVDF, (b) PLLA/PVDF (50/50), (c) PLLA/PVDF/L-0-2-8(50/50/3), (d) PLLA/PVDF/C-1-2-7(S-2400)(50/50/3), (e) PLLA/PVDF/C-1-2-7(S-4800)(50/50/3) and (f) PLLA/PVDF/C-1-2-7(S-6300)(50/50/3)

Figure 9. FT-IR spectra for stretched (a) PLLA, (b) PVDF, (c) PLLA/PVDF/L-0-2-8(50/50/3), (d) PLLA/PVDF/C-1-2-7(S-2400)(50/50/3), (e) PLLA/PVDF/C-1-2-7(S-4800)(50/50/3) and (f) PLLA/PVDF/C-1-2-7(S-6300)(50/50/3)


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Table 3. The relationship between β-phase content and draw ratio Sample Draw ratio β-phase content PVDF 2.5 70 % PLLA/PVDF 0.05 0 (50/50) PLLA/PVDF/L-0-2-8 0.9 45% (50/50/3) PLLA/PVDF/C-1-1-8(S-2400) 3.0 75% (50/50/3) PLLA/PVDF/C-1-2-7(S-2400) 3.6 83% (50/50/3) PLLA/PVDF/C-1-2-7(S-4800) 4.9 95% (50/50/3) PLLA/PVDF/C-1-2-7(S-6300) 3.9 90% (50/50/3) Discussion Although there existed some discrepancies in the compatibilization efficiencies of RC polymers with different molecular structures, all of them were significantly higher than that of RL polymers. TEM investigations on the microstructure of the PLLA/PVDF blends further showed that the interfacial area of the blends compatibilized by RC polymers was slightly thicker than that by RL polymers (Figure 10) and this was in accordance with the rheological properties that G’ of RC-compatibilized blend was higher than that by RL, because of the more contribution from the interface. A significant difference between the blends compatibilized by RL and RC was that a few white particles with a diameter of about 50 nm existed in the PVDF phase for RL polymer (Figure 10 and Figure S4).[59] During reactive blending, the epoxide groups of GMA reacted with the end groups of PLLA and the PLLA chains were thus grafted onto the backbone of PMMA (Scheme S1). As shown in Figure 11, for RL polymer, the PMMA backbone resided in the PVDF dispersed phase and the grafted PLLA chains extended into and entangled with the PLLA matrix. We speculated that as the PLLA grafting reaction proceeding, the average molecular weight between two grafting sites in the PMMA backbone decreased. When this value was approaching the entanglement molecular weight between PMMA and PVDF (Me12=5700),[42] the interaction between the PVDF phase and the PMMA backbone became weakening and the PLLA grafted RL polymers tended to be expelled from the interface and pulled into the PVDF phase to form micelles (indicated by arrows in Figure 10) or pulled out to the PLLA matrix. As shown in Figure 11, for PLLA grafted RC polymers, the molecular structure was totally different from that by RL



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polymers, because they were more symmetric and had two kinds of side chains. One of them was PMMA, which was derived from the MMA macromer, the other was PLLA from grafting reaction. The PMMA and PLLA side chains penetrated into the PVDF and PLLA phase, respectively. This in some extent could counterbalance the weakening of the interaction between the PMMA backbone and the PVDF phase, incurred by the grafting reaction. Thus it was the appropriate number and length of the PMMA side chains that increased the stability of the RC polymers at the interface and efficiently lowered the interfacial tension and inhibited coalescence of the dispersed phase. At the same time, the PMMA side chains were also found to affect the crystallization behaviors of PVDF phase in the RC-compatibilized blends. Unlike the neat PVDF which had well-developed spherulites, the PMMA side chain might induce the formation of some imperfect crystallites (i.e. needle-like), which was prone to orient or rotate along the stress field. The energy was consumed during this process and thus the mechanical properties of the blends were more or less improved after compatibilization.[23]

Figure 10. High-magnification TEM images of PLLA/PVDF (50/50) blends compatibilized by 3wt% of (A) L-0-2-8; (B) C-1-2-7 (S-2400); (C) C-1-2-7 (S-4800); (D) C-1-2-7 (S-6300). The scale bar was 200nm


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Figure 11. Schematic diagrams PLLA grafted RC or RL polymers at the interface of PLLA/PVDF (matrix/dispersed phase). The micelles formed by pull-in have a core of PLLA (color yellow) and a shell of PMMA (color blue); those by pull-out have an inverse phase (PMMA core and PLLA shell) Conclusion A series of RL and RC polymers with different architectures have been successfully synthesized. Firstly, we found that compared to our former report, as the GMA content increased from 10 wt% to 20 wt% in both the RL and RC polymers, the compatibilization efficiencies were significantly improved. Secondly, the RC polymers were better compatibilizers than their linear counterparts, especially those with moderate length of side chains. We ascribe the superiorities of RC polymers to the existence of side chains, which in some extent could counterbalance the weakening of the interaction between PMMA backbone and PVDF phase, as the average molecular weight between two grafting sites decreased with grafting reaction proceeding. The PMMA side chains were also found to affect the crystallization behaviors of PVDF phase. The reactive compatibilization endowed the PLLA/PVDF blends with excellent toughness. The blends compatibilized by RC polymers, especially the ones with 20 wt% of GMA and PMMA side chains with a Mn of 4800 or 6300 g mol-1 (i.e. C-1-2-7(S-4800) or C-1-2-7(S-6300)), could be stretched to a high draw ratio (more than about 4.0) at room temperature, during which process, the α phase of PVDF was almost completely transformed into the piezoelectric β phase. The thus obtained PLLA/PVDF blends with the sustainable PLLA as the matrix had good adhesion with other materials like metals and cells and would be a perfect candidate in energy harvesting devices.



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SUPPORTING INFORMATION AVAILABLE The characterization of PMMA-COOH and MMA macromer, 1H-NMR spetrum of RC polymer, SEM and TEMimages of PLLA/PVDF blends compatibilized by RL or RC polymers, stress-strain curves, XRD and FT-IR spectra of neat components and PLLA/PVDF blends and scheme of the reaction between PLLA and RC polymers. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Authors *Email: [email protected]. TEL: 86-571-2886-7206; FAX: 86-571-2886-7899. Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT This work was financially supported by the National Natural Science Foundation of China (21244009, 21304026, 51173036, 21374027) and Program for New Century Excellent Talents in University(NCET-13-0762).The authors wish to express their thanks to editors and the reviewers for their valuable and professional comments and suggestions.

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Table of Contents Use Only Dramatic Improvement in Toughness of PLLA/PVDF Blends: the Effect of Compatibilizer Architectures

Wenyong Dong, Hengti Wang, Fanglu Ren, Junqing Zhang, Meifeng He, Tao Wu, Yongjin Li*

Synopsis: Both the adhesion and toughness of PVDF were dramatically improved by blending with PLLA and compatibilized by RC polymers.


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Dramatic Improvement in Toughness of PLLA ... - ACS Publications

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