In Situ Formation of Microfibrillar Crystalline Superstructure: Achieving

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In Situ Formation of Microfibrillar Crystalline Superstructure: Achieving High-Performance Polylactide Chunhai Li, Ting Jiang, Jianfeng Wang, Hong Wu, Shaoyun Guo, Xi Zhang, Jiang Li, Jiabin Shen, Rong Chen, and Ying Xiong ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b06705 • Publication Date (Web): 14 Jul 2017 Downloaded from http://pubs.acs.org on July 18, 2017

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ACS Applied Materials & Interfaces

In Situ Formation of Microfibrillar Crystalline Superstructure: Achieving High-Performance Polylactide

Chunhai Lia, Ting Jiangb, Jianfeng Wanga, Hong Wua*, Shaoyun Guo, Xi Zhang, Jiang Lia*, Jiabin Shena, Rong Chena, Ying Xionga

a

The State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of

Sichuan University, Chengdu 610065, China b

School of Chemistry, State Key Laboratory of Biotherapy of Sichuan University, Chengdu 610065,

China

ABSTRACT：As one biobased and biodegradable polyester, polylactide (PLA) is widely applied in disposable products, biomedical devices and textile. Nevertheless, due to its inherent brittleness and inferior strength, simultaneously reinforcing and toughening of PLA without sacrificing its biodegradability is highly desirable. In this work, a robust assembly consisting of compact and well-ordered microfibrillar crystalline superstructure (FCS) surrounded by slightly oriented amorphism, is achieved by a combined external force field. Unlike the classic crystalline superstructure such as shish-kebabs, cylindrites and lamellae, the newfound FCS with diameter of about 100 nm and length of several tens of micrometers is aggregated with well-aligned crystalline

* To whom correspondence should be addresses. (Prof. Wu, Email: [email protected], Fax: 86-028-85466077) * To whom correspondence should be addresses. (Prof. Li, E-mail [email protected], Fax: 86-028-85466077)

ACS Paragon Plus Environment


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nano-fibers. FCS can serve as discontinuous fiber to self-reinforce the amorphous PLA, more importantly, FCS can also act as rivets to pin the propagating fibrillar crazes leading to the formation of dense fibrillar crazes during stretching, which dissipates so much energy and translates the failure of PLA from brittle to ductile. Consequently, PLA with FCS exhibits exceptionally simultaneous enhancement in ductility, strength and stiffness, outperforming normal PLA with increments of 728%, 55% and 70% in elongation at break, strength, and modulus, respectively. Therefore, FSC exhibits competitive advantages in achieving high-performance PLA even for other semicrystalline polymers. More significantly, this newfound crystalline superstructure (FCS) provides a new structural model to establish the correlation between structure and performance. KEYWORDS：crystalline fibrils, polylactide, flow-induced crystallization, crazes, toughness, performance

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INTRODUCTION

Poly(lactide) (PLA), a biobased and biodegradable aliphatic thermoplastic polyester, is one of the most promising alternatives to petrochemical-derived plastics.1,2 Nevertheless, PLA should be modified to enable its applications more widespread due to its inherent brittleness, inferior strength and stiffness.3,4 However, traditional approaches to modify PLA are prominently hampered by strength-ductility trade-off dilemma, i.e., large sacrifice in strength and stiffness after blending with flexible components 5-10 and poor ductility and toughness after adding rigid fillers.11,12 Considering that PLA as a semicrystalline polymer, its physical properties are governed by


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its crystalline architecture.13-15 As a result, regulating its crystalline architecture and the secondary components simultaneously enables PLA based blends or composites with intrinsic structural basis for high-performance. And this can be exemplified by the large increase of toughness in PLA/PCL(Poly (ε-Caprolactone)) blends containing PLA matrix with high crystallinity,16 the unexpected combination in strength, stiffness and ductility (56.4 MPa, 1702 MPa, and 92.4%) permitted by precise control of PBS (Poly (Butylene Succinate)) or PBAT (Poly (Butylene Adipate-Co-Terephthalate)) phase morphology as well as PLA crystalline architecture,17-19 and the superior strength, stiffness and ductility (119.4 MPa, 2342 MPa, and 8.7%) conferred by aligning the rigid fillers like ramie fibers20 or ZnO whiskers21 to induce the epitaxial growth of PLA lamellae. However, homogeneity is paramount to PLA as a biomaterial because any additive composed of different chemicals might lead to negative effects on its transparency, biocompatibility and biodegradability.22,23 From this perspective, achieving high-performance PLA without any additive is of great significance. According to the theoretical calculation, polymers along chain direction possess tensile strength and Young’s modulus up to 30 and 250 GPa, respectively.24,25 Thus, PLA can be extremely strengthen and stiffen through full orientation of PLA chains, which is usually achieved via solid-state extrusion,26,27 uniaxial hot-drawing,28,29 and fiber spinning.30-32 Unfortunately, PLA with highly oriented chains generally exhibits much poor toughness and ductility due to the limitation of molecular mobility. Given this, rather than try to align all molecular chains, only implanting well-oriented PLA fibers or tapes into un-oriented PLA matrix endows PLA with unexpected balance in strength, stiffness and ductility (77.5 MPa, 3200 MPa



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and 65.2%).33-35 In fact, natural load-bearing materials such as nacre, spider silk, tendon and wood are well known for their superior combination in strength, stiffness and toughness. Those materials are characterized by their hierarchic structures, which is consisted of compact and well-ordered reinforcing elements surrounded by soft and amorphous energy-absorbing matrix.36,37 Naturally, it stimulates us to achieve high-performance PLA by regulating its hierarchic structure, wherein the compact and highly oriented crystalline architecture serves as the reinforcing element while the soft and un-oriented amorphism serves as energy-absorbing matrix. To date, flow-induced crystallization has been verified as the most efficient way to regulate crystalline architecture of semicrystalline polymers.38 However, due to the semi-rigid molecular backbone and low molecular weight of PLA, regulating its crystalline architecture such as crystalline morphology and orientation needs much intense flow filed,39 which offers only by some elaborately designed processing methods. By far, with oscillation shear injection molding, Li et al. first fabricated neat PLA with shish-kebabs concentrating on the skin of the injection bars14. With layer multiplying co-extrusion, our recent work reported a kind of compact shish-kebabs.15 Specially, with the twostep melt spinning, i.e., PLA melt firstly undergoes extensional and shearing field provided by the spinneret and then followed by the elongational field provided by the take-up device; PLA fibers consisting of microfibrils was achieved, endowing PLA with unprecedented balance in strength, stiffness and ductility (350 MPa, 3200 MPa, 40%),31 suggesting that fibrillar microstructure is an ideal reinforcing element. Nevertheless, the formation of fibrillar microstructure in PLA requires extremely intense extensional and shearing field, which is realized only by fiber spinning at


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present.40 Therefore, it is still a challenge to introduce fibrillar microstructure from one dimensional PLA fiber to three dimensional PLA articles. With this challenge in mind, in present work, a number of layer multiplying elements (LMEs), designed containing tow fishtail channels, are connected in series to superpose the extensional and shearing field created by those fishtail channels. By connecting the LMEs in series, the intensity of extensional and shearing field could increase exponentially (Scheme 1(a)). By far, the intense extensional and shearing field provided by LMEs has been verified to be effective in in-situ fibrillation of the dispersed phase,41,42 the orientation of high-aspect ratio fillers.43-45 As aforementioned, the main purpose here is to strengthen and toughen PLA simultaneously by structuring a designed “assembly”. Hence, similar to the two-step melt spinning,30,31 the extensional and shearing field offered by LMEs together with the elongational filed provide by the take-up device was imposed on the PLA melt (Scheme 1). Interestingly, PLA with a new microfibrillar crystalline superstructure (FCS) was successfully achieved. The hierarchic structure of FCS and the corresponding strengthening and toughening mechanism for PLA were investigated systematically. As FSC can simultaneously strengthen and toughen PLA, FSC exhibits competitive advantages in achieving high-performance PLA or even other semicrystalline polymers.

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EXPERIMENTAL SECTION Materials. Commercially available PLLA, comprising 2% DLA (trade name 4032D), was

purchased from Nature Works Co. (U.S.A.), whose weight- and number-average molecular



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weights were 2.23×105 and 1.06×105 g mol−1, respectively. It has a density of 1.24 g/cm3, and its MFR is 7 g/10min at 210 ℃, 2.16Kg. Sample Preparation. PLLA 4032D was dried in a vacuum oven at 80 ℃ for 12h before extrusion, and the dried pellets were extruded through multistage stretching extrusion system. As illustrated in Scheme 1(a), the system contains a single-screw extruder, layer multiplying elements (LMEs) and a take-up device. The LMEs were designed containing tow fishtail channels to provide extensional and shearing field for polymer melt. As shown in the bottom of Scheme 1(a), in one of the LMEs, the polymer melt was firstly sliced into two sections by a divider, and then each section flowed through a fishtail channel; before flowing into next LME, the polymer melt was vertically recombined. As shown in Scheme 1(b), by connecting the 9 LMEs in series, both the converging ratio along thickness direction (x, z plane) and the extending ratio along planar direction (x, z plane) can reach 512, and the volume deformation ratio can reach 512×512, of the order 105. Meanwhile, PLLA without LME was also applied to produce the control samples. The temperatures from the hopper to the LMEs are 160, 185, 180 and 170 ℃, respectively. Specially, 170 ℃, around the melting point of PLLA, is set for LMEs for increasing the relaxation time of PLLA chains. The feeding rate was 23 g/min. The drawing speed of the take-up device was fixed at 7.2, 16.2 and 35.6 mm/s, which the corresponding draw ratio is calculated to be 67%, 155% and 280%, respectively. The as-prepared samples were named as PLLA-x-y-T, the denotation x refers to the number of LME, y refers to the draw ratio, and T means the extruded sheets which were annealed at 95 °C for 2.5 h.


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Scheme 1. (a) schematic of multistage stretching extrusion system; (b) the converging and extending ratio of PLLA melt after flow through 9 LMEs; (c) structure evolution for the crystalline fibrils.

Scanning Electronic Microscopy (SEM). The extruded sheet was placed in liquid nitrogen for 2h, and then the surface of the fractured sample blocks were chemically etched by a watermethanol (1:2, v:v) solution containing 0.025 mol/L sodium hydroxide for 12h at 30 °C. After that, the etched surface was cleaned by distilled water and ultrasonication. A field-emission SEM (JSM5900LV, Japan) was utilized with the accelerated voltage of 10 kV, all samples were sputter-coated with gold prior to observation. Two-Dimensional Small-Angle X-ray Scattering (2D-SAXS). 2D-SAXS measurements were performed with the Xeuss 2.0 system of Xenocs, France, with a sample-to-detector distance of 2500 mm providing scattering vector s (s = (2sin θ)/λ) range from 0.008 to 0.183 nm-1. A multilayer focused Cu Kα X-ray source (GeniX3D Cu ULD, Xenocs SA, France, λ = 0.154 nm)



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and scatterless collimating slits were used during the experiments. The size of the primary X-ray beam at the sample position was 0.5 ×0.5 mm2. SAXS images were recorded with a Pilatus 300K detector of Dectris, Swiss (680 pixels × 600 pixels, pixel size = 172 μm). A silver behenate standard was used to calibrate the parameters of the scattering geometry (i.e., beam center and sample-to-detector distance). The 1-D and 2-D projection method46,47 and Ruland rule48,49 were used to explore the crystalline structure. Two-Dimensional Wide-Angle X-ray Diffraction (2D-WAXD). 2D-WAXD was also conducted on the Xeuss 2.0 system of Xenocs, with a sample-to-detector distance of 172 mm. The silver behenate standard was used to calibrate the parameters of the scattering geometry. The overall crystallinity, 𝑋𝑐 , was calculated according to the following equation: Xc =

∑ Acryst ∑ Acryst +∑ Aamorp

(1)

where, 𝐴𝑐𝑟𝑦𝑠𝑡 and 𝐴𝑎𝑚𝑜𝑟𝑝 are the fitted areas of crystal and amorphous region, respectively. The orientation of PLLA crystals can be estimated using Hermans’ orientation factor (𝑓𝐻 ). For a set of ℎ𝑘𝑙 planes, Herman’s orientation function is defined as follows: 𝑓𝐻 =

3⟨𝑐𝑜𝑠2 𝜑⟩𝑙𝑘ℎ −1

(2)

2

where 𝑐𝑜𝑠 2 𝜑 is the orientation factor defined as 𝜋/2

⟨𝑐𝑜𝑠 2 𝜑⟩𝑙𝑘ℎ =

∫0

𝐼(𝜑) 𝑐𝑜𝑠2 𝜑 𝑠𝑖𝑛𝜑 𝑑𝜑

𝜋/2

∫0

𝐼(𝜑) 𝑠𝑖𝑛𝜑 𝑑𝜑

(3)

where 𝜑 is the azimuthal angle and 𝐼(𝜑) is the scattered intensity along the angle 𝜑 . The azimuthal intensity distribution 𝐼(𝜑) was analyzed at 2θ = 16.7 °, where the peak represents the (200, 110)α reflections in PLLA.50 The mean size of crystal domain 𝐿ℎ𝑘𝑙 of extruded sheet before


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and after thermal treatment were calculated by Scherrer equation.51 Differential Scanning Calorimeter (DSC). A DSC Q2000 (TA Instruments, U.S.A.) was used to study the thermal features such as melting and crystallization behaviors of the PLLA. The PLLA samples around 8 mg were heated from 40 to 210 °C at a heating rate of 10 °C/min under nitrogen atmosphere. The crystallinity (𝑋𝑐 ) can be calculated by subtracting the enthalpy of cold crystallization from the enthalpy of melting by using following equation: 𝑋𝑐 =

∆𝐻𝑚 −∆𝐻𝑐𝑐 0 ∆𝐻𝑚

× 100%

(4)

where ∆𝐻𝑚 is the enthalpy of melting for PLLA, ∆𝐻𝑐𝑐 is the enthalpy of cold crystallization, 0 and ∆𝐻𝑚 is the enthalpy of melting for a 100% crystalline of PLLA (93.7 J/g).52

Polarized FTIR Test. The molecular orientation of different PLLA samples was measured by the Thermo Nicolet iS10 FTIR spectrometer with a resolution of 2 cm-1 and an accumulation of 32 scans. The test slice samples with thickness about 20 μm were cut by a rotary Microtome (YD-2508B) along the flow direction from the center of virgin and thermal treatment samples, respectively. In order to quantitatively characterize the molecular orientation of PLLA, FTIR spectra were also recorded in the transmittance mode when the polarizer was rotated. The dichroic ratio D and Herman orientation function 𝑓 can be obtained using the following equations:53 𝐴

𝐷 = 𝐴‖

(5)

⏊

𝐷−1

𝑓 = 𝐷+2

(6)

where A‖ and A⏊ are absorbance intensities parallel and perpendicular to the flow direction, respectively. For PLLA, the band at 921 cm-1 belongs to the coupling of C-C backbone stretching linked to the CH3 rocking mode, which is considered as the exclusively crystal-sensitive band of



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α-crystals, while the band at 955 cm-1 is related to the amorphous phase of PLLA.54 From bands at 921 cm-1 and 955 cm-1, both the orientation function of the crystalline phase 𝑓𝑐 and amorphous phase 𝑓𝑎 can be determined. Mechanical Testing. According to GB/T 1040-92, standard tensile test using dumbbell shaped sample was conducted on the instrument (model CMT-4104) with a crosshead speed of 50 mm/min at ambient temperature. Each measurement was repeated at least six times in the same conditions, and the average values were presented with standard deviation.

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RESULTS AND DISCUSSION

Superior Mechanical Performances of PLLA with Microfibrillar Crystalline Superstructure (FCS). Mechanical performances, especially strength and ductility, are crucial for the applications of PLLA as structural products. As shown in Figure 1(upper), the mechanical performances of PLLA with 9 LMEs (Figure 1(a) and (b)) are absolutely superior to the performances of PLLA without LME (Figure 1(c) and (d)), PLLA with 9 LMEs show high yield stress as well as distinct ductility. Detailed tensile properties regarding yield strength, Young’s modulus and elongation at break are summarized in the Figure 1(bottom). Appreciable enhancement of strength and ductility for PLLA with 9 LMEs is quite clear, with unexpected promotion of elongation at break and yield strength up to 35.6% and 90.8 MPa in PLLA-9-280%, and 23.3% and 119.1 MPa in PLLA-9155%-T (after annealing) compared to the initial values of 4.5% and 63.3 MPa for control sample (PLLA-0-67%). Additionally, as shown in Figure S1, PLLA regardless of draw ratio exhibit well optical clarity. It should be noted that the clarity of PLLA with 9 LMEs is ascribed to their


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Figure 1. (a~d): typical stress-strain curves of various samples (upper); (e) detailed mechanical results regarding yield strength, Young’s modulus, and elongation at break (bottom).

nano-size crystal, while the clarity of PLLA with 0 LME is ascribed to their low crystallinity, which will be discussed in next section. In summary, with the intense extensional and shearing field provided by LMEs and the elongational flow provided by the take-up device, the strong, tough, stiff and transparent PLLA was successfully achieved. Besides, to clearly assess the contribution of present work, the performances of PLLA obtained by other self-reinforcing methods were



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summarized in Table S1. Microfibrillar Crystalline Superstructure (FCS). The SEM was first performed to visually exam the crystalline superstructure of PLLA. Surprisingly, rather than the classic shish-kebabs or cylindrites, in the samples with 9 LMEs, it is compact FCS aligning along flow direction (Figure 2(A1~B3)). Such FCS strongly depends on the draw ratio, only the precursor of FCS is formed with 67% draw ratio (Figure 2(A1)), whereas FCS orienting with their long axis along the flow direction is successfully achieved with the draw ratio above 155% (Figure 2(A2) and 2(A3)).

Figure 2. SEM micrographs of (A1): PLLA-9-67%; (A2): PLLA-9-155%; (A3): PLLA-9-280%; (B1): PLLA9-67%-T; (B2): PLLA-9-155%-T; (B3): PLLA-9-280%-T; (C1): PLLA-0-67%-T; (C2): PLLA-0-155%-T; (C3): PLLA-0-280%-T; T means the sample was annealed at 95 ℃ for 2.5 h; the blue arrow indicates the flow direction.


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After annealing at 95 ℃ for 2.5h, FCS displays a much compact array (Figure 2(B1~B3)), and this may result from the epitaxial growth of crystal on the surface of pristine FCS. As for the samples obtained without LME, regardless of the draw ratio, only random PLLA lamellae are visible (Figure 2(C1~C3)). Besides, a skin-core structure is expected here because: (i) the flow rate of PLLA melt decreases parabolically from the center to wall of the fishtail channel;55 and (ii) more time for the crystallization of PLLA at the center of the extrudate due to its slower cooling rate. As illustrated in Figure 3, the oriented FCS is focused in the core section of the extrudate. Very close to the skin, i.e., 5 μm away from the surface, only the random crystalline entities are observed as shown in Figure 3(a). At the distance of 70 μm away from the skin, FCS occurs sporadically (Figure 3(b)). Further increasing the distance, PLLA crystallizes into well-aligned FCS (Figure 3(c), 3(d), 3(e)). Moreover, compared with the FCS in out-layer (Figure 3(a), 3(b), 3(c)), it seems that the FCS in core-layer (Figure 3(d), 3(e)) shows smaller diameter and much compact array.

Figure 3. SEM micrographs showing the different regions of the annealed PLLA with 9 LMEs and 280% draw ratio (PLLA-9-280%-T), the value inset in the up-right indicates the distance from the skin of the extruded sheet.



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Orientation Structure of PLLA with FCS. SAXS and WAXD measurements were further performed to gain more information of FCS. As shown in Figure 4(A1~A3), the spot reflections in their 2D-WAXD patterns show the highly oriented α crystal in PLLA with 9 LMEs, regardless of the draw ratio. The reflection of lattice plane (200, 110)α locating on the equator is strongest, while the reflection of (203)α and (206)α locating on quadrant are comparatively weak.50 The weak reflections of (203)α, (206)α and the absence of (015)α indicate that the FCS consists of imperfect crystalline structure, which can be also clearly verified by the 1D-WAXD diffraction profiles

Figure 4. 2D-SAXS scattering and 2D-WAXD patterns of PLLA before annealing: (A1), (A2) and (A3), WAXD patterns for PLLA-9-67%, PLLA-9-155% and PLLA-9-280%, respectively; (B1), (B2) and (B3), WAXD patterns for PLLA-0-67%, PLLA-0-155% and PLLA-0-280%, respectively; (a1), (a2) and (a3), SAXS scattering for PLLA-9-67%, PLLA-9-155% and PLLA-9-280%, respectively; (b1), (b2) and (b3), SAXS scattering for PLLA-0-67%, PLLA-0-155% and PLLA-0-280% , respectively; (flow direction is vertical).


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(Figure S2). In order to identify the amount of misorientation of the FCS, Herman’s orientation function, 𝑓𝐻 was calculated according to equation 2. As shown in Table 1, only with draw ratio 67%, the 𝑓𝐻 reaches to 0.9 and even reaches to 0.96 when the draw ratio increases to 280%, suggesting the highly oriented crystals in PLLA with 9 LMEs. After annealing, 𝑓𝐻 decreases slightly, indicating the less orientation of the new formed crystal. As to PLLA without LME, only amorphous structure manifested by their amorphous halo in 2D-WAXD (Figure 4(B1~B3)) were obtained. Interestingly, as for the SAXS scattering for PLLA with 9 LMEs, streak scattering along the equator is displayed (Figure 4(a1~a3)). The streak scattering indicate that PLLA with 9 LMEs show the firbillar superstructure, no trace of lamellar structure, which is consistent with the crystalline superstructure revealed by SEM (Figure 2 and 3). Actually, the streak scattering in SAXS is usually observed in polymer fibers and its interpretation is more complicated since it may contain several contributions including the fibrillar superstructure, deformed voids and the surface reflection/scattering of the fibers.30,56 However, unlike fibers, the sample here is a block, so the streak scattering cannot be originated from the surface reflection/scattering. Besides, PLLA with 0 LEM present no detectable SAXS scattering (Figure 4(b1~b3)), indicating the draw process provided by the take-up device does not produce the deformed voids. Hence similar to the Kevlar fibers,56,57 the streak scattering in PLLA with 9 LMEs (Figure 4(a1~a3)) is originated from its crystalline fibrillar structure, not from the void. As for the annealing samples, similar but a bit different 2D-WAXD patterns and 2D-SAXS scattering were obtained because of the evolution of crystal structure during annealing (Figure S3).



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Figure 5. first heating curves of PLLA with 9 LMEs and 0 LME, respectively; (a) before annealing; (b) after annealing at 95 ℃ for 2.5 h.

DSC was also employed to provide direct insights into the crystal information of PLLA with 9 LMEs and 0 LEM, respectively. The resultant DSC heating curves for PLLA are plotted in Figure 5. As shown in Figure 5(a), two transitions, i.e., cold crystallization and melting endotherm are successively displayed on the heating curves. PLLA with 9 LMEs regardless of draw ratio present approximate cold crystallization enthalpy (about 10.3 J/g) and crystallinity (about 35.6%). Nevertheless, PLLA without LME display both the big cold crystallization enthalpy (about 29.3 J/g) and low crystallinity (below 7%). More interestingly, the glass transition temperature shifts from about 59.8 ℃ in PLLA without LME to about 67.8 ℃ in PLLA with 9 LMEs, indicating that the chains in PLLA with 9 LMEs become well packed and less mobile due to their high crystallinity and molecular orientation. After annealing (Figure 5(b)), the cold crystallization peak vanishes and all samples present approximate crystallinity. However, the glass transition temperature of PLLA with 9 LMEs is still higher than that of PLLA without LME, which may result from the highly oriented crystals whose molecular arrangement are more extended.


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Figure 6. Polarized FTIR spectroscopy of PLLA with 9 LMEs: (A) PLLA-9-280%, (B) PLLA-9-280%-T (after annealing); (a1~d1) polar diagrams of amorphous absorbance peaks at 955 cm-1 for PLLA-9-67%, PLLA-9-155%, PLLA-9-280% and PLLA-9-280%-T, respectively; (a2~d2) polar diagrams of crystalline peaks at 921 cm-1 for PLLA-9-67%, PLLA-9-155%, PLLA-9-280% and PLLA-9-280%-T, respectively.

PLLA as a semicrystalline polymer, the physical properties of PLLA strongly depend on its crystallites as well as amorphism. Consequently, polarized FTIR was further employed to explore the orientation not only in crystalline region but also in amorphism. Figure 6(A) shows the absorbance of polarized FTIR of PLLA-9-280% while Figure 6(B) shows PLLA-9-280%-T (after annealing). The polarization angle means the angle between the flow direction and polarized direction of the polarizer. Specifically, the band at 921 cm-1 belonging to the crystalline 103 helix



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is assigned to the coupling of C-C backbone stretching linked to the CH3 rocking mode (E mode, i.e., its maximum would appear perpendicularly to the oriented direction of PLLA chains), while the band at 955 cm-1 is related to the amorphism of PLLA (A mode, i.e., its maximum would appear parallelly to the oriented direction of PLLA chains).54,58 As shown in Figure 6(A) and 6(B), both the bands at 921 and 955 cm-1 display variation along polarization angle, indicating neither the crystalline region nor the amorphous region is isotropic. In detail, the absorbance of 921 and 955 cm-1 as a function of the polarization angle are plotted as a polar diagram (Figure 6(a1~d2)). As shown in Figure 6(a1~d1), the maximum and minimum absorbance of the band at 955 cm-1 is located at the 0 and 90 polarization angle, respectively, which indicates the molecules in amorphism orientate along the flow direction. On the contrary, as shown in Figure 6(a2~d2), the maximum and minimum absorbance of the band at 921 cm-1 is located at the 90 and 0 polarization angle, respectively. As the crystalline band at 921cm-1 is assigned to the E mode, its maximum would appear perpendicularly to the oriented direction of PLLA molecules. Consequently, the molecules in crystalline region also orientate along the flow direction. The difference value between maximum and minimum represents the degree of orientation, and the orientation function, 𝑓, calculated by equation 6, is shown in Table 1. The 𝑓 of the amorphism (955 cm-1) for all samples are below 0.11, indicating that the molecules in amorphism only slightly oriented. In contrast, the 𝑓 of the crystalline region (921 cm-1) for all samples regardless of draw ratio approach -0.5, indicating that the molecules in crystalline region almost align along the flow direction. In summary, based on above structural investigation and mechanical performances, it can be concluded that high-performance PLLA has been achieved in the present work by


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structuring a designed “assembly”, wherein compact and well-ordered FCS serve as an ideal reinforcing element, while the slightly oriented amorphism serves as soft and amorphous energyabsorbing matrix. Quantitative Analysis of the Hierarchic Structure of FCS. The average length of the structural units of FCS (𝐿𝑓 ) and their alignment along the flow direction (𝐵𝜑 ) can be determined using the method proposed by Perret and Ruland.48,49 This method consists of determining the integral breadth of a series of azimuthal profiles taken at different values of s along the equator. The integral breadth (𝐵𝑜𝑏𝑠 ) is defined according to the equation: 𝜋⁄2

1

𝐵𝑜𝑏𝑠 (𝑠) = 𝐼(𝑠,𝜋⁄2) ∫−𝜋⁄2 𝐼(𝑠, 𝜑)𝑑𝜑

(7)

where φ is the azimuthal angle and I(s,φ) is the background corrected intensity. If Lorentzian profile can be used to describe the orientation distribution, 𝐵𝑜𝑏𝑠 is then related to 𝐵𝜑 and 𝐿𝑓 by 𝐵𝑜𝑏𝑠 (𝑠) = 𝐵𝜑 + 𝐿

1 𝑓𝑠

(𝐶𝑎𝑢𝑐ℎ𝑦 − 𝐶𝑎𝑢𝑐ℎ𝑦)

(8)

All results in this study are well fitted with Lorentz functions as demonstrated in the inset of Figure 7. On the basis of above equation, 𝐿𝑓 and 𝐵𝜑 can be obtained from a plot of 𝐵𝑜𝑏𝑠 versus 1/s (Figure 7), and the resulting parameters were summarized in Table 1. The magnitude of 𝐿𝑓 first increases from the 300.2 nm for the draw ratio 67% to 437.8 nm for the draw ratio 155%, and then drops to the 322.5 nm for draw ratio 280%. We assume that 𝐿𝑓 has a distribution and the measured value represents a mean value. It is conceivable that some structural units of FCS may break because they probably carry the most loading when the samples were taken-up with the draw ratio 280%, which leads to the decrease of 𝐿𝑓 . The 𝐵𝜑 only slightly decreases from 16.5o for the draw ratio 67% to 15.2o for the draw ratio 155%, nevertheless, it sharply drops to 8.2o as the draw ratio



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further increase to 280%. After annealing, both the 𝐿𝑓 and 𝐵𝜑 increase slightly.

Figure 7. Plot of azimuthal integral breadth (𝐵𝑜𝑏𝑠 ) as a function of 1/s showing the Ruland method used to separate out the contributions of the misorientation of (𝐵𝜑 ) and the length of structural units of FCS (𝐿𝑓 ) at low scattering vectors (s). Inset shows a typical background corrected azimuthal intensity profile fitted with a Lorentz function.

The crystal size, i.e., the coherent scattering region dimension in the direction perpendicular to the diffracting planes, is also calculated based on the Scherrer Equation.51 The lattice plane (203) is chosen as the reference plane since the overlapped (210), (110) is not suitable for Scherrer Equation. Moreover, the diameter of the structure units is conversed from the crystal size (𝐿203 ) based on the geometric crystallography (equation S8, Supporting Information). As shown in Table 1, the draw ratio does not affect the diameters, i.e., all samples present similar diameters, about 17.7 nm. After annealing, this value slightly increases about to 19.7 nm. Both the longitudinal and traverse structure of FCS were also estimated through the 1D and 2D SAXS projections established by Stribeck,46,47 more detailed information is available in


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Supporting Information. The long period along flow direction can be calculated using the Bragg equation, L=1/s*, here L is long period, and s* represents the position of the maximum of shoulder peak in the 1D projection curves (Figure S5(a)). As shown in Table 1, the long period along the axis of FCS is about 15.2 nm, which is much smaller than that reported in PLLA with perfect crystal (22.8 nm),14 further indicates that FCS consists of imperfect crystal. After annealing, the long period slightly increase to about 19.6 nm. As for the 2D projection curve (Figure S5(b)), its projection intensity drop sharply with increasing S12 and no peak can be distinguished, which indicates the irregular scatters along the traverse direction of the flow direction. It is very worthy that both the length ( 𝐿𝑓 , from 300.2 to 437.8 nm) and the diameter (about 17.7 nm) of the structural units calculated based on X-ray data are noticeably smaller than the magnitude of the FCS shown in SEM (Figure 2, 3). Before SEM observations, the samples were sputter-coated with about 8 nm thick gold, which may cause the deviation of SEM micrograph from its original crystalline morphology. Nevertheless, combing the morphology observed from SEM with the analysis of the X-ray data, the hierarchic structure of FCS can be speculated. As shown in Scheme 2(C), the X-ray scattering coherent region, i.e., the crystalline nano-fibers or the structural units, with the length 𝐿𝑓 (from 300.2 to 437.8 nm), diameter (about 17.7 nm) and long period (about 15.2 nm) are bundled to form FCS (Scheme 2(A) and (B)). In particular, as shown in Scheme 2(C), rather than alternating crystalline/amorphous structure, it is alternating crystalline/semicrystalline structure that consists of the crystalline nano-fibers. As shown in table 1, slight difference in the crystallinity calculated from WAXD and DSC is observed for the annealed samples, and this does make sense because of their different theoretical principles.



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However, for the samples before annealing, the crystallinity calculated based on WAXD is far below that based on DSC. Consequently, it can be concluded that the crystalline nano-fibers is consisted of imperfect crystals, which are not regular enough to scatter the X-ray well while are regular enough to display endotherm during melting. Moreover, the 1-D projected intensity curves (Figure S5(a)) only show a faint shoulder peak, suggesting the weak difference in electron density along the fiber axis of FCS (along 𝑠3 ). This also suggests the poor period along the fiber axis of FCS. Therefore, the alternating crystalline/semicrystalline structures is proposed as the structural model of the crystalline nanofibers, because if it is the alternating crystalline/amorphous structures, the difference in electron density along the fiber axis of FCS should be stronger and a stronger peak should be also discovered in the 1-D projected intensity curves. And the crystalline structures here indicate the PLLA chains crystallize at regular arrangement, while the semicrystalline Table 1. Crystalline parameters of PLLA with FCS. Crystallinity Herman Orientation Functiona Parameters of Crystalline Nano-Fibers Sample

XRD DSC

XRD

Polarized IR

Lengthb

Long Period

d Orientationb Diameterc along Fibril Axis

(%)

(%) (200),(110) 955 cm-1 921 cm-1

(nm)

(degree)

(nm)

(nm)

PLLA-9-67%

14.1

35.6

0.90

0.11

-0.48

300.2

16.5

17.7

14.6

PLLA-9-155%

14.5

33.1

0.95

0.07

-0.49

437.8

15.2

17.7

15.2

PLLA-9-280%

15.2

37.9

0.96

0.06

-0.49

322.5

8.2

17.9

17.8

PLLA-9-67%-T

49.5

48.3

0.82

0.08

-0.31

319.0

22.6

19.7

19.6

PLLA-9-155%-T 49.7

47.6

0.85

0.06

-0.36

463.5

18.1

19.2

18.8

PLLA-9-280%-T 49.6

46.1

0.88

0.06

-0.38

342.6

10.8

19.9

20.4

Data are calculated with a: Hermans’ orientation equation; b: Ruland rule; c: Scherrer equation and d: Projection methods.


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structures here indicate the PLLA chains crystallize at moderate regular arrangement. Besides, it should also be noted that the structural difference between the FCS and the fibrillar microstructure in PLLA fibers. The FCS is bundled with crystalline nano-fibers while the fibrillar microstructure found in PLLA fibers is stacked with alternating crystalline/amorphous structure.30,31 Mechanism for Strong, Tough, and Stiff PLLA with FCS. SAXS and WAXD investigations of PLLA-9-280% during stretching provide a pathway to explore the structural evolution and thereafter ascertain the corresponding strengthening and toughening mechanism. As shown in Figure 8, the SAXS scattering and WAXD patterns of pristine extrudate (ε = 0%) indicate the FCS presents fibrillar morphology. Cross-shaped SAXS scattering originating from the fibrillar crazes (Figure 9(A)) can be observed when the stress was applied. The SAXS model for fibrillar

Figure 8. Typical stress-strain curve of the PLLA-9-280% along with two-dimensional SAXS scattering (left) and WAXD patterns (right) of PLLA-9-280% drawn to ε=0%, 4.4%, 8.6%, 16.0%, 22.9%, 29.3%, 35.6% (drawn rate is 50mm min-1 and the draw direction is vertically).



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craze has demonstrated that the streak scattering parallel to the loading direction (meridional direction) is ascribed to the total reflection at fibrillar crazes/intact PLLA interfaces, while the two equatorial lobes perpendicular to the loading direction (equatorial direction) are caused by the alternating fibrils/micro-voids structure (Figure 9(A)) in the fibrillar crazes.59,60 As the situation here is that the equatorial scattering of stressed sample is a superposition of the equatorial lobes scattering caused by the alternating fibrils/micro-voids structure (Figure 9(A)) and the streak scattering originated from the FCS (Figure 4(a3)), it is difficult to estimate the parameters of fibrillar crazes with SAXS craze model.48, 61 However, as shown in Figure 9(A), the alternating fibrils/micro-voids can be treated as two-dimensional lattice with quasi long-range order. Both the mean diameter of fibrils (d) and inter-fibrillar distance (D), can be derived from the onedimensional correlation function proposed by Strobl.62,63 The electron density correlation function K(z) is derived from the inverse Fourier transformation of the experimentally intensity distribution I(s) as follows: ∞

𝐾(𝑧) =

∫0 𝐼(𝑠12 )𝑐𝑜𝑠(𝑠12 𝑧)𝑑𝑠12 ∞

∫0 𝐼(𝑠12 )𝑑𝑠12

(9)

where z is perpendicular to the loading direction and 𝐼(𝑠12 ) is obtained by integrating along equator. The resultant values are summarized in Table 2, the diameter of the fibrils (d) in fibrillar craze first reaches its maximum at 7.5 nm when the sample was stressed to yield point and then drops slightly during the following stretching, while the inter-fibrillar distance (D) keeps stable at about 21.1 nm during the whole stretching. Moreover, as shown in Figure 8, fibrillar crazes generate cross-shaped SAXS scattering ( ∆𝐼𝑐𝑟𝑎𝑧𝑒 ) concentrating on equatorial (∆𝐼𝑐𝑟𝑎𝑧𝑒‖ ) and meridional

(∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ) sections.48, 61 Hence the propagation of fibrillar crazes can be assessed by


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Figure 9. (A) schematic of fibrillar craze; (B) one-dimensional correlation function curves of PLLA-9-280 drawn to different stretching ratio, the upper-left inset presents the method of determining diameter of fibrils (d) and inter-fibrillar distance (D), respectively; (C) WAXD profiles and (D) DSC heating curves for PLLA-9-280 drawn to different stretching ratio.

Table 2．Parameters of fibril-crazes and crystalline evolution of PLLA-9-280 during stretching. Parameters of Fibrillar-Craze

Crystallinity

Strain(ε) (%)

Diameter of Fibrils (d) (nm)

Inter-Fibrillar Distance (D) (nm)

Meridional Ratio ∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ∆𝐼𝑐𝑟𝑎𝑧𝑒 + ∆𝐼𝑐𝑟𝑦𝑠𝑡

XRD (%)

DSC (%)

Diameter of Nano-fibers (nm)

0

--

--

0

14.1

37.9

17.9

4.4

7.5

21.1

0.42

12.4

36.8

16.2

8.6

7.1

21.1

0.75

12.2

36.3

16.0

16.0

6.9

21.3

0.81

12.5

36.1

16.0

22.9

7.0

20.5

0.81

12.3

36.0

16.3

29.3

7.0

21.3

0.84

12.3

36.1

16.1

35.6

6.9

21.5

0.77

11.1

35.4

13.2



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tracking the variation of meridional ratio, ∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ⁄∆𝐼𝑐𝑟𝑎𝑧𝑒 . Additionally, the scattering originated from fibrillar crazes (∆𝐼𝑐𝑟𝑎𝑧𝑒 ) is much stronger than that from the FCS (∆𝐼𝑐𝑟𝑦𝑠𝑡 . only concentrating on equational section) because the much higher electron difference between the fibrils and micro-voids in fibrillar crazes (Figure 9(A)) compared to that in PLLA between crystalline

and

amorphous

region.

Therefore,

the

meridional

ratio,

∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ⁄∆𝐼𝑐𝑟𝑎𝑧𝑒 ≈ ∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ⁄(∆𝐼𝑐𝑟𝑎𝑧𝑒 + ∆𝐼𝑐𝑟𝑦𝑠𝑡 ) and the ∆𝐼𝑐𝑟𝑎𝑧𝑒 + ∆𝐼𝑐𝑟𝑦𝑠𝑡 can be obtained by integrating the whole scattering region. As shown in Table 2, the meridional ratio, ∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ⁄∆𝐼𝑐𝑟𝑎𝑧𝑒 , increases quickly in the case of draw ratio below 8.6% and then keeps stable during the further stretching, after break, its magnitude drop slightly. Besides, the digital photos taken after break (Figure S6) indicate that there is no necking in dumbbell-shaped samples, only the stress whitening is observed. Therefore, the continuous increase of volume of the deforming PLLA should be ascribed to the micro-void created by forming fibrillar crazes continuously.64 In other words, the fibrillar crazes were formed continuously over the whole stretching region. However, as both the diameter of the fibrils (d) and the meridional ratio (∆𝐼𝑐𝑟𝑎𝑧𝑒⏊ ⁄∆𝐼𝑐𝑟𝑎𝑧𝑒 ) keep stable during stress plateau and the successive stress hardening region, it can be speculated that the propagating fibrillar crazes are terminated as soon as they grow to a fixed size. And the most possibility is that the fibrillar crazes are terminated as soon as they propagate to the FCS. As for the crystalline evolution during stretching, the crystallinity based on WAXD (Figure 9(C)) and DSC (Figure 9(D)), and the diameter of crystalline nano-fiber (Scheme 2(B) and (C)) decrease slightly when the sample was stressed. This suggests that some crystalline structure is broke into amorphism during the stretching.


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Crazing involves a large degree of plastic deformation localized within very thin zones and their presence has been linked with enhanced toughness in polymers such as high-impact polystyrene (HIPS), where crazes are known to absorb up to 90% of the strain energy.65 However, their presence has also been linked to the onset of brittle fracture, which leads to catastrophic failure through the transition from crazes to crack. Fortunately, the case here is that fibrillar crazes were terminated as soon as they grow to a fixed size, no trace of the transition from crazes to crack. Based on the above analysis, the mechanism for the strong, tough, and stiff PLLA with FCS is proposed. As shown in Scheme 2(A), (B) and (C), PLLA with FCS can be seen as self-reinforcing system where the discontinuous FCS reinforces the amorphous PLLA matrix. It is well-known that the mechanical performances of fiber composites largely depend on stress transfer between fiber and matrix, as the same chemical structure of FCS and amorphous PLLA, a tenacious interphase

Scheme 2. (A), (B) and (C) schematic of the hierarchic structure of microfibrillar crystalline superstructure (FCS); (D) termination of fibrillar crazes by FCS; (E) transition from fibrillar crazes to cracks in normal PLLA.



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is expected. Therefore, PLLA with FCS exhibits high tensile stress and stiffness (modulus). Pioneering researches illustrated that amorphous PLLA is a brittle material with a failure resulting from the quick transition of crazes to crack.49, 66 As shown in Scheme 2(E), the amorphous PLLA exhibits poor ductility because of the catastrophic break caused by the quick transition from fibrillar crazes to crack. Nevertheless, as shown in Scheme 2(D), for PLLA with FCS, when the initial fibrillar crazes propagate to FCS, FCS serves as rivets to pin fibrillar crazes and then terminate the fibrillar crazes. Moreover, as the FCS stacks with well parallel and compact array, the fibrillar crazes can be efficiently terminated and thereafter denser fibrillar crazes rather than crack are formed (Scheme 2(C)). The formation of fibrillar crazes rapidly releases the stress concentrating at local scale, and the formation of denser fibrillar crazes can overall avoid the break caused by stress concentration. Most importantly, the generation of denser fibrillar crazes dissipates great energy. Therefore, PLLA with FCS presents well ductility. 

CONCLUSION

A designed assembly of a natural load-bearing material, consisting of compact and well-ordered self-reinforcing elements (microfibrillar crystalline superstructure, FCS) surrounded by slight oriented amorphism, is achieved by combining intense extensional and shearing field and elongational flow through an extrusion technology. Without any additive and sacrifice of biodegradability, such compact and well-ordered FCS endows PLLA with exceptional strength and stiffness for: (i) FCS serves discontinuous fiber to self-reinforce the amorphous PLLA matrix; and (ii) stress can be well transferred from the amorphous PLLA matrix to FCS due to the tenacious interphase generated from their same chemical structure. Most importantly, FCS also confers


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PLLA with unprecedented ductility, with an increment of 728% in elongation at break compared with the normal PLLA. Structural analysis shows that FCS serves as rivets to pin fibrillar crazes as soon as the fibrillar crazes propagate to the FCS. Moreover, as FCS embeds in amorphous PLLA with well-oriented and compact array, denser crazes can be formed during stretching, which dissipates great energy and prolongs the catastrophic break caused by the transition of fibrillar crazes to crack. Therefore, PLLA with FCS exhibits unprecedented ductility. Supporting Information. 1D-WARD diffraction profiles, 2D-SAXS scattering and 2D-WAXD patterns of PLLA after thermal treatment at 95 ℃, digital photos taken before and after the samples were stretched, and a table for the comparison of mechanical performance of PLLA with different fabrication methods. Acknowledgements Financial support of the National Natural Science Foundation of China (51573118 and 51227802), Program for New Century Excellent Talents in University (NCET-13-0392), Sichuan Province Youth Science Fund (2015JQ0015) and the Program for Changjiang Scholars and Innovative Research Team in University (IRT-15R48) are gratefully acknowledged. 

REFERENCE

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(3) Spinella, S.; Cai, J.; Samuel, C.; Zhu, J.; McCallum, S. A.; Habibi, Y.; Raquez, J.-M.; Dubois, P.; Gross, R. A. Polylactide/Poly (ω-hydroxytetradecanoic acid) Reactive Blending: A Green Renewable Approach to Improving Polylactide Properties. Biomacromolecules 2015, 16, 18181826. (4) 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, 2899-2916. (5) Dong, W.; Wang, H.; Ren, F.; Zhang, J.; He, M.; Wu, T.; Li, Y. Dramatic Improvement in Toughness of PLLA/PVDF Blends: the Effect of Compatibilizer Architectures. ACS Sustainable Chem. Eng. 2016, 4, 4480-4489. (6) Li, T.; Zhang, J.; Schneiderman, D. K.; Francis, L. F.; Bates, F. S. Toughening Glassy Poly (lactide) with Block Copolymer Micelles. ACS Macro Lett. 2016, 5, 359-364. (7) 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, 11203-11214. (8) Dhar, P.; Kumar, A.; Katiyar, V. Magnetic Cellulose Nanocrystal Based Anisotropic Polylactic Acid Nanocomposite Films: Influence on Electrical, Magnetic, Thermal, and Mechanical Properties. ACS Appl. Mater. Interfaces 2016, 8, 18393-18409. (9) Ojijo, V.; Sinha Ray, S.; Sadiku, R. Role of Specific Interfacial Area in Controlling Properties of Immiscible Blends of Biodegradable Polylactide and Poly [(Butylene Succinate)-co-Adipate]. ACS Appl. Mater. Interfaces 2012, 4, 6690-6701.


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