Natural

Article pubs.acs.org/JPCB

Toward Stronger Transcrystalline Layers in Poly(L‑lactic acid)/Natural Fiber Biocomposites with the Aid of an Accelerator of Chain Mobility Huan Xu, Lan Xie, Xin Jiang, Xu-Juan Li, Yue Li, Zi-Jing Zhang, Gan-Ji Zhong,* and Zhong-Ming Li* College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Chengdu, 610065, Sichuan, People’s Republic of China ABSTRACT: Formation of transcrystalline layer probably enhances the interfacial adhesion of poly(L-lactic acid) (PLLA)/natural fiber biocomposites as confirmed by this work. We found that a crystallization accelerator, poly(ethylene glycol) (PEG), improved chain mobility of PLLA and thus enhanced the growth kinetics of ramie fiber-induced transcrystallinity (TC). The direct observation of polarized optical microscopy during isothermal crystallization revealed that large-sized TC with rapid growth was produced after adding PEG. It could be exemplified by the case at 125 °C that the growth rate of TC developed in PLLA10 (containing 10 wt % PEG) achieved 6.1 μm/min, which was nearly triple that of pure PLLA (2.1 μm/min). And interestingly enough, spherulitic nucleation proceeding was largely restricted because it was difficult to fulfill the critical size for stable nuclei due to the increased chain mobility. Meanwhile, combining the effective nucleation activity of ramie fibers and acceleration virtue of PEG offered the chance to form prevailing TC texture, instead of rich spherulites dominated in pure PLLA. The local structure (including lamellar structure and molecular orientation) of transcrystalline layers was further determined, which indicated that TC presented α crystal form and random lamellar packing derived from the moderate nucleating ability. To our surprise, the single fiber reinforced composite samples containing prevailing TC textures achieved remarkably higher strength compared to that of pure PLLA samples with poorly developed transcrystalline layers, as demonstrated by the single-fiber pull-out test.

■

INTRODUCTION The past decades have witnessed the intensely growing research interest of poly(L-lactic acid) (PLLA) as an environmental friendly and versatile biopolymer that is already available from renewable resources and allows for biotical or abiotical degradation in the environment.1−3 Additionally, PLLA is highly attractive in various applications, including fibers, the packaging industry, biological and medical usage like bone tissue engineering, and articles for daily use such as cups and bottles.4−9 Nevertheless, the intrinsic brittleness, low ductility, and low resistance to gas and water make pure PLLA unable to be directly used in the above-mentioned areas.10−14 Aiming at optimizing PLLA properties, ongoing attempts are made to modify its intrinsic brittleness.15−18 Directly blending with various plasticizers, such as tributyl cirate,19 dibutyl sebacate,20 poly(propylene glycol)21 and poly(ethylene glycol) (PEG),22,23 has been demonstrated to be a promising and convenient gateway in terms of considerably enhanced chain mobility of PLLA.24 For example, PEG has been found to be one of most effective accelerators for improving chain mobility, and the elongation at break of PLLA can climb up to 300%.23,24 However, the direct usage of plasticizers seems to sacrifice the stiffness of PLLA to some extent, which is primarily reflected in the dramatically decreased strength and modulus, as well as unsatisfied impact toughness.15,19,25 Hence simultaneously reinforcing the plasticized PLLA probably represents a compelling approach to complement © XXXX American Chemical Society

the sacrificial strength and modulus. Among the various reinforcing agents, natural fibers which reflect their distinct advantages in terms of low density, easy accessibility, renewability and biodegradability, and prominent property profiles have to be the outstanding candidates.26−29 Recent developments involving biocomposites of PLLA and natural fibers such as jute, flax, sisal, and ramie fibers, are rapidly appearing and further strengthening the possibility of substituting the so-called commodity polymers with biopolymers.30−33 Such biocomposites are highly appreciated as highperformance materials that can absolutely degrade.33,34 We recently proved that ramie fiber was indeed a favorable reinforcing element for PLLA; therein the tensile strength and modulus of the biocomposites gradually increased with the fiber content, surprisingly achieving 91.3 and 2977 MPa upon the addition of 30 wt % ramie fiber from initial 65.6 and 1468 MPa for pure PLLA, respectively, without compromising toughness and ductility.16,35 Reasonably, it is expected a promising balance of mechanical performance in the ternary biocomposites containing both natural fibers and plasticizer. As for a polymeric composite reinforced with fiber, the interfacial structure between fiber and polymer matrix presents a straightforward influence on the mechanical properties of Received: September 9, 2013 Revised: November 16, 2013

A

dx.doi.org/10.1021/jp409021q | J. Phys. Chem. B XXXX, XXX, XXX−XXX

The Journal of Physical Chemistry B

Article

composite materials.36 As an interesting, columnar interfacial morphology, transcrystallinity (TC) texture probably forms in the vicinity of foreign fibers. From the practically technological point of view, TC appears to show its superiority in controlling the interfacial properties of various composites.37−39 Intrigued by this potential and with the goal of creating considerable interfacial adhesion for high-performance composites, numerous researchers set out to realize the underlying role of TC.39−41 On the basis of elaborate exploration, it is of great interest to find that TC show preferred crystallite orientation relative to the fiber,42 enhanced fiber-matrix adhesion, reduced stress concentration and cavitation at the fiber ends,39 a stronger and stiffer layer in the direction of shear than the bulk phase,40 a protective sheath around fibers,43 and a powerful mechanical interlock created by filling the pores of the rough fiber surface with this layer.44 Thereby the local stress can be transferred from the matrix across the interphase in a more efficiently favorable manner and, finally, resulting in preferable composites with superior performance. Unfortunately, few attempts have been made to reveal the hierarchical structure of TC of PLLA induced by natural fibers, and the intrinsic properties of PLLA TC, and favorable conditions for transcrystallization remain still poorly understood, especially in a ternary PLLA biocomposite system containing natural fibers and chain accelerator.45 It naturally arouses our curiosities: What is the structure at different scales (e.g., molecular orientation, lamellar superstructure, and crystalline structure) of PLLA TC induced by natural fibers, how does the accelerator affect the formation and growth of TC for the modified PLLA, and what is the potential relationship between the formed TC and mechanical performance? Inspired by the above unclear, but interesting questions, we focus on the interfacial morphology of PLLA/ramie fiber biocomposites with or without a chain accelerator of PLLA in this work. Water-soluble and highly flexible PEG with short chain length is used here because it shows considerable compatibility with PLLA and excellent efficiency in improving crystallization kinetics of PLLA.20,46 Polarized optical microscopy (POM) and scanning electron microscopy (SEM) were employed to observe the crystallization kinetics and lamellar structure of PLLA TC, while scanning microbeam two-dimensional wideangle X-ray diffraction (2D-WAXD) analyses with high spatial resolution were carried out to examine the crystalline structure and molecular orientation in TC. Furthermore, interfacial shear strength was evaluated through single-fiber pull-out test to understand the relationships between the interphase and interfacial properties.

was blended with various contents of PEG (0, 3, 5, and 10 wt %) in dichloromethane (CH2Cl2), and then a thin film with a thickness of approximately 10 μm was generated by solution casting for subsequent characterization or testing. For the sake of briefness, the plasticized films are hereafter referred to as PLLA3, PLLA5, and PLLA10 samples, which contain 3, 5, and 10 wt % PEG, respectively. POM Observation. To explore the formation and growth of TC of PLLA, a single ramie fiber with a diameter of ca. 17 μm was sandwiched between two pieces of the aforementioned thin film and melted by a Linkam CSS450 hot stage at 180 °C for 3 min. The sample was then cooled rapidly to a series of preset temperatures (i.e., 125, 130, and 135 °C), at which the isothermal crystallization would proceed for 30 min. The temperature protocol is schematically presented in Figure 1A.

Figure 1. (A) Temperature protocol for fiber-induced transcrystallization, Tc stands for the temperature preset for isothermal crystallization (i.e., 125, 130, and 135 °C). (B) Schematic for characterization of crystalline structure of spherulite and TC. This digital image shows the crystalline morphology of the sample PLLA10 after isothermal crystallization for 30 min at 125 °C.

At this stage, an Olympus BX51 polarizing optical microscopy (Olympus Co., Tokyo, Japan) equipped with a Micro Publisher 3.3 RTV CCD was applied to trace the morphology evolution during crystallization. It is noteworthy that the optical micrographs presented in this paper were all taken under crossed polarizers. Scanning Microbeam 2D-WAXD Measurement. In order to determine the detail crystalline morphology of spherulite and TC formed in the biocomposites, scanning microbeam 2D-WAXD measurements were performed for the sandwiched films after isothermal crystallization on the hot stage. 2D-WAXD measurements were conducted at the beamline BL15U1 of the Shanghai Synchrotron Radiation Facility (SSRF, Shanghai, China). We scanned a sample with the microfocused X-ray beam from one side to another side of the sample across the ramie fiber with a step of 20 μm as diagramed in Figure 1B. The monochromated X-ray beam with a wavelength of 0.124 nm was focused to an area of 3 × 2.7 μm2 (length × width), and the distance from sample to detector was held at 147.5 mm. The 2D-WAXD images were collected with an X-ray CCD detector (Model SX165, a resolution of 2048 × 2048 pixels, Rayonix Co. Ltd., America). Additionally, the crystallinity was calculated by the ratio of the area under the resolved Gaussian crystalline peaks to the total area under the unresolved diffraction curve.47 SEM Observation. SEM observation was performed aiming to investigate the lamellar structure of spherulite and TC directly. The sandwiched films after isothermal crystallization on the hot stage were etched in a water−methanol (1:2 by volume) solution containing 0.025 mol/L of sodium hydroxide

■

EXPERIMENTAL SECTION Materials. Commercially available PLLA comprising around 2% D-LA was purchased from Nature Works (trade name 4032D), and its weight-average molecular weight and numberaverage molecular weight were 2.23 × 105 g/mol and 1.06 × 105 g/mol, respectively. PEG with a nominal weight-average molecular weight of 3.35 × 103 g/mol was obtained from Dow Chemical Company under the trade name Carbowax. Ramie fibers used were bast fibers from the outer culm of the ramie plant with a density of 1.5 g/cm3 and an average diameter of 17 μm ranging from 10 to 40 μm, which were kindly supplied by Yuzhu Plant Fiber Industrial Co. Ltd., Sichuan, China. Sample Preparation. PLLA and ramie fibers were first dried at 100 °C under vacuum overnight, in order to avoid hydrolysis degradation during isothermal crystallization. PLLA B



Article

Figure 2. POM observations for crystallization morphology of isothermal crystallization at 125 °C of (a) pure PLLA, (b) PLLA3, (c) PLLA5, and (d) PLLA10.

for 11 h at 25 °C; subsequently, the etched films were cleaned by using distilled water and ultrasonication. A field-emission SEM (Inspect F, FEI, Finland) was utilized to carefully examine the etched surfaces, which were sputter-coated with gold prior to SEM observations, and the accelerated voltage was kept at 5 kV. Single-Fiber Pull-Out Test. The crystallized PLLA, PLLA3, PLLA5, and PLLA10 samples containing single ramie fiber-induced TC after isothermal crystallization at the preset temperatures were immediately removed from the hot stage and cryogenically quenched in liquid nitrogen, while the amorphous single-fiber composites were prepared to make the control samples by quenching the molten composites in liquid nitrogen directly. A piece of sheet (10 × 1 mm, length × width) was carefully prepared from the crystallized or amorphous samples. The single ramie fiber is perpendicular to the length direction of the sheet, and the length of the fiber embedded in the matrix is 1 mm. To perform the pull-out test, a plastic block with a connecting groove in the midcourt line was machined, and the sheet was adhered to the block with the fiber located on top of the groove, while the fiber outside the matrix was glued to a paper card.48 Then, partially embedded fibers were pulled out from the matrix using an Instron universal test instrument (Model 5576, Instron Instruments, USA) with a crosshead speed of 1 mm/min and a gauge length of 20 mm. At least ten replicates were tested for each type of sample, while the average value was reported.

growing front of spherulites in the bulk. Ultimately, transcrystallization is ceased as it just reaches the radius of approximately 30 μm. At the time point of 30 min, spherulites further grow to almost cover the bulk wherein the transcrystalline layer is restricted in a limited space. In clear contrast, the addition of PEG tends to turn the tables. Figure 2b−d reveals that, the observed TC textures obtain much larger sizes in the PLLA3, PLLA5, and PLLA10 samples compared to that of the pure PLLA sample, therewith develop to the prevailing crystalline structure with the aid of a small amount of PEG. The size of TC appears to increase with the PEG content. Specifically, the transcrystalline layer formed in PLLA3 reaches a dramatically elevated value of around 70 μm, which further climbs up to 110 μm for PLLA5 and arrives at 160 μm for PLLA10. Compared to the thin transcrystalline layer and numerous spherulites in the pure PLLA system, the thickness of TC texture of PLLA10 is tremendously promoted over 5 times. The generation and domination of large-sized TC primarily lies in the preferential nucleation at the fiber surface and subsequently the largely enhanced growth of TC, which are attributed to the effective nucleation activity of ramie fibers35 and significantly accelerated crystallization kinetics of TC stimulated by PEG,49 respectively. Ultimately, only sporadically dispersed spherulites form due to the poor nucleation ability of bulk PLLA, which provides more available space for the sufficient growth of TC without excessive disturbance of spherulitic proceeding. As for the heterogeneous nucleation of ramie fiber, it was reported that unique topography on the natural fiber surface like small-scale grooves and valleys tend to cause the thermal stress concentration and thus enhance the heterogeneous nucleation, while probable interactions between cellulose macromolecules of natural fibers and molecular chains of the polymer matrix close to the fiber and various impurities adhered to the fiber surface are also assumed to be responsible for the effective nucleating ability of natural fibers.44,45 The intrinsic heterogeneous nucleation activity can be applied to improve crystallinity for the biopolymers with poor crystallization kinetics.35

■

RESULTS AND DISCUSSION TC Morphology of PLLA Induced by Ramie Fiber. Figures 2−4 present POM observations of crystalline morphology developing gradually in the pure PLLA, PLLA3, PLLA5, and PLLA10 samples during isothermal crystallization. For the crystallization temperature set at 125 °C as illustrated in Figure 2a, the pure PLLA shows strong nucleation at the ramie fiber surface to form TC in the appearance of columnar layer, and quite a few spherulitic nuclei are generated in the bulk matrix, simultaneously. The growth of transcrystalline layer is prematurely hindered by the impingement of the C



Article

Figure 3. POM observations for crystallization morphology of isothermal crystallization at 130 °C of (a) pure PLLA, (b) PLLA3, (c) PLLA5, and (d) PLLA10.

Figure 4. POM observations for morphology evolution of isothermal crystallization at 135 °C of (a) pure PLLA, (b) PLLA3, (c) PLLA5, and (d) PLLA10.

Compared to the case of pure PLLA generated compact spherulites, the observation of remarkably decreased density of spherulites after adding PEG arouses our interest to understand the effect of PEG on the formation of PLLA spherulites. The common point is that the formation of polymeric spherulite is essentially dependent on nucleation and crystal growth.50 PLLA chains seem to form stable nuclei easily at relatively low crystallization temperature of 125 °C, since lots of spherulites can be observed. Thus, the crystallization of pure PLLA at 125 °C is controlled by crystal growth. The addition of PEG seems to change this crystallization model, making a transition from growth-controlled crystallization to nucleation-controlled crystallization, since we observed that spherulite is hard to form in a PLA/PEG system and nucleation needs be accomplished by the ramie fiber. With the addition of PEG as a chain accelerator, it is difficult to fulfill the critical size of spherulitic nuclei because

the chain mobility dramatically augments due to miscible PEG chains. As such, given the same crystallization temperature and time, only scarce spherulites form in the PLLA/PEG systems, whereas transcrystallization is preferentially induced by ramie fiber-aided nucleation. This circumstance leaves sufficient space for the development of TC enclosing the ramie fiber. It is thus clear that the evolution of spherulitic nucleation is significantly controlled by accelerating chain mobility of PLLA with the aid of PEG, and a nucleation-controlled crystallization tends to process with the presence of PEG. As the crystallization temperature increases to 130 °C, the nucleation density of spherulite in the bulk decreases considerably in all the four samples because it is somewhat difficult for heterogeneous nucleation to proceed at higher temperature, while ramie fibers can effectively stimulate heterogeneous nucleation in the interface, and therefore D



Article

Figure 5. Quantitative analysis on time dependence of radius of TC of (a) pure PLLA, (b) PLLA3, (c) PLLA5, and (d) PLLA10 during isothermal crystallization at (A) 125 °C, (B) 130 °C, and (C) 135 °C, and on (D) growth rates of TC of the four samples at different temperatures.

transcrystalline layers form and develop along the fibers in the pure and plasticized PLLA samples, as evidently shown in Figure 3. Likewise, transcrystalline layers with large sizes develop in the PLLA3, PLLA5, and PLLA10 samples, which mainly lies in less spatial hindrance and fast growth after introducing PEG, achieving considerable radius of around 80, 150, and 140 μm, respectively. However, the thickness of TC produced in pure PLLA remains at a low level of about 40 μm, attesting to the fact again that TC growth is spatially restricted by the fierce competition with adjacent spherulites. The isothermal crystallization is further performed at 135 °C for the pure PLLA and modified samples, the crystalline morphologies are illustrated in Figure 4. The nucleation in the bulk and heterogeneous nucleation at fiber surface display slightly different appearance. Nucleation for spherulite in the bulk is further restricted, principally indicating that polymer chains are difficult to fulfill the critical size for stable nucleus at this present temperature. On the other hand, although ramie fibers still can effectively induce TC texture, the heterogeneous nucleation seems to be weak and only a few nuclei were observed in the vicinity of fiber surface. Of great interest, Figure 4d reveals the PLLA10 sample just develops transcrystalline layer without any spherulites. In the case of high crystallization temperature accompanied by the addition of PEG, PLLA chains possess extremely high mobility, which substantially weakens the ability of nucleation for spherulite. It desirably meets the viewpoint proposed by Folkes and Moon, in which they demonstrated the simultaneous formation and growth of spherulites should be avoided during the transcrystallization.40,51 Low strain to failure, low failure energy, and reduced interfacial strength in the transcrystalline composite materials were observed when large, well-developed spherulites coexisted in the matrix. Additionally, their assumption was proposed to be the absence of interconnecting material between crystallites, thus leading to the fast propagation of cracks in these areas.40,51 As previously found, it is difficult to isolate the contribution of TC from the crystallites in the bulk, resulting in the properties of the transcrystalline layer itself are lacking investigation. Our work suggests TC texture can be separately fostered by

adjusting the temperature and PEG content wherein the intrinsic properties of TC can be desirably investigated. To acquire quantitative analysis regarding the growth of TC, we measured the radius of TC versus the crystallization time as shown in the curves of Figure 5A−C. These linear curves suggest that TC regularly presents progressive growth at a constant rate. Distinctly, the growth curves of TC in all PLLA/ PEG system lie above those of pure PLLA, which means the size of TC is remarkably enlarged by adding the accelerator of chain mobility. According to the slopes of the growth curves, values of growth rates of TC are produced, and the growth rate as a function of crystallization temperatures is summarized in Figure 5D. Generally, the growth rates are prone to decrease as the crystallization temperature rises, especially for the PLLA10 sample. Figure 5D reveals that, TC of PLLA10 grows at a rate of ∼6.1 μm/min at 125 °C, whereas the growth rate sharply drops down to approximately ∼5.2 and ∼3.5 μm/min at the crystallization temperatures at 130 and 135 °C, respectively. Compared to crystallization temperature, the addition of PEG within 3−10 wt % gives rise to a more significant influence on the growth rate of TC. For instant, the growth rates of TC in pure PLLA present the lowest level of approximately 2.1 μm/ min, whether at low or high crystallization temperatures. In the PLLA/PEG system, a dramatic rise in growth rates of TC is evidently achieved. Additionally, the growth rate is proportional to the PEG content for PLLA3 and PLLA5 samples, jumping to the higher levels of around 3.3 and 6.0 μm/min, respectively, with slight variations for various crystallization temperatures. No more increase is, interestingly, observed as long as the PEG content exceeds 5 wt %, the PLLA10 sample fostered at 125 °C obtains comparable value of growth rate like the case of PLLA5. Note that, although experiencing a gradual decline of TC growth rate at 130 and 135 °C, TC growth of PLLA10 is still maintained at a relatively high level compared to that of the PLLA3 sample. In view of the preferential generation and fast growth of TC by controlling the PEG addition, we probably unlock the potential to manufacture the natural fiber reinforced biopolymer biocomposites containing plenty of large-sized TC, E



Article

Figure 6. A representative series of scanning microbeam 2D-WAXD patterns of the typical sample containing a transcrystalline layer (PLLA5) after isothermal crystallization for 30 min at 130 °C.

Figure 7. Selected 2D-WAXD patterns of spherulite and TC of (a) pure PLLA, (b) PLLA3, (c) PLLA5, and (d) PLLA10 after isothermal crystallization for 30 min at 125 °C, 130 °C, and 135 °C.

from which we can expect the enhanced interfacial properties and thus the improved bulk performance. Although a great quantity of experimental evidence usually based on differential scanning calorimeter results have obtained a clear description of the significant effects of plasticizers with low molecular weight such as biocompatible PEG and poly(propylene glycol), in enhancing the crystallization kinetics for PLLA,21−23 unfortunately, little attention is paid to facilitate the transcrystallization of PLLA in the natural fiber reinforced biocomposites. Herein, we confirmed the synergistic effects of PEG and ramie fiber in controlling the crystalline morphology, which originates from the heterogeneous nucleation activity of ramie fibers, improved growth rates of TC and suppressed nucleation of spherulites after introducing PEG. As a

consequence, it renders the rapid development of TC superstructure, which cannot be attained by using ramie fibers alone. Noticing the significant facilitation of transcrystallization and suppression of the nucleation for spherulite, we attempt to obtain a further understanding in terms of the role of PEG. During the isothermal crystallization, PEG is kept in the molten state at the preset crystallization temperatures far above the melting point of PEG (∼53 °C),23 and the PEG chains with relatively short length and high mobility are miscible with amorphous PLLA.24 Once the heterogeneous nucleation is stimulated at the fiber surface, for which it is believed that transcrystallization proceeds easily thanks to the high nucleation efficiency of ramie fiber,35 PEG will aid PLLA molecules in folding into lamellae more rapidly and form TC F



Article

Figure 8. 1D-WAXD intensity profiles of spherulite and TC for the four samples after isothermal crystallization for 30 min at 125 °C, 130 °C, and 135 °C.

with large radius in an energetically favorable manner compared to the pure PLLA system. On the other hand, phase separation probably occurs as a result of isothermal crystallization of the PLLA/PEG samples, which likely disturbs the spherulitic nucleation in the bulk and transcrystallization at the fiber surface to some extent. Thus sporadic spherulites and slow growth of TC are evidently shown in Figures 2d, 3d, 4d, and Figure 5D, and this phenomenon becomes more evident at higher temperature probably due to more severe phase separation caused by higher mobility of PEG chains. Local Crystalline Structure of PLLA TC Induced by Ramie Fiber. There are comprehensive articles in determine the crystalline structure of TC by employing wide-angle X-ray diffraction.52 Unfortunately, commonly used X-ray facility has a large beam size reaching hundreds (or thousands) of micrometers, which cannot distinguish TC and spherulites due to poor spatial resolution. To separately identify the fine structure of TC and spherulites, we made use of microbeam 2D-WAXD, which can focus to an area of 3 × 2.7 μm2 to characterize the fine structure of transcrystalline layers. Figure 6 illustrates a series of representative 2D-WAXD patterns for a typical sample containing the transcrystalline layer as the X-ray microbeam scans from one side to another side across the ramie fiber. Apparently, Figure 6a,b demonstrates the existence of amorphous area due to the presence of diffuse halo pattern, and subsequently spherulites are detected as identified by the reflections of lattice planes (200)/(110) and (203) of PLLA crystals, which is clearly depicted in Figure 6c,d. As the sample moves step by step, another new crystal reflection corresponding to the lattice plane (200) of cellulose I crystal in ramie fiber

emerges as shown in Figure 6e,f, which reminds us of the transcrystalline layer at the ramie fiber. Afterward, the microbeam step-by-step leaves the ramie fiber. Figure 6g−l suggests that the reflection patterns vary from strong crystalline diffraction rings to halo amorphous diffraction, and then evolve to moderate crystalline diffraction, suggesting that these areas are mixtures of crystallites and amorphous area. Typical diffraction patterns of spherulite and TC of pure PLLA and modified PLLA samples fostered at various temperatures are selected out for a clear comparison of their crystalline structure as envisaged in Figure 7. What is immediately noticeable from Figure 7 is that there is no obvious difference of crystal orientation and diffraction circles between spherulite and TC for all samples, except for the additional diffraction arcs from highly oriented crystallites in ramie fiber, i.e., random orientation of PLLA TC is detected at or around the ramie fiber. The results are not in line with the previous observation in some other polymer systems, in which a main viewpoint holds that TC is well developed and organized from a more compact crystal packing and probably from a preferred crystalline alignment.53,54 Note that, a preferential lamellae growth along the b-axis of PLLA crystal in the transcrystalline layer was obtained by placing PLLA films melted between two poly(tetrafluoroethylene) sheets, which may indicate that different subtract also affect the lamellar orientation of TC.55 For poly(tetrafluoroethylene) with strong heterogeneous nucleating ability, preferential orientation could be induced. The random lamellar packing of PLLA TC in the vicinity of ramie fiber surface essentially results from the moderate nucleating capability of ramie fibers. We further G



Article

manner as those of spherulite due to the moderate nucleating ability of ramie fiber. Relationships between the Interphase Structure and Interfacial Properties. It is recognized that the structure of interphase is of great significance to the adhesion between reinforcing fiber and matrix in a composite.58−61 The existence of transcrystalline layer is a favorable feature to improve interfacial adhesion and provide a powerful mechanical interlock created by filling the pores of the rough fiber surface with this layer, especially for the composites with poor matrixfiber interfacial adhesion.44 However, the properties of the transcrystalline layer itself lack investigation, especially for TC formed in PLLA biocomposites, which is due to the difficulty of isolation of the contribution of TC from the crystallites in the bulk. Herein we evaluated the interfacial properties for the single-fiber composites containing TC of various thicknesses through the microbond shear testing, and the typical microbond pull-out curves for ramie fibers partially embedded in pure PLLA and PLLA/PEG films are plotted in Figure 10. The

attempt to investigate the crystal form of TC on the basis of WAXD patterns, from which the 1D-WAXD intensity profiles are extracted as presented in Figure 8. For specific comparison, we note that the diffraction peaks of spherulite and TC reveal the well coincidence for all curves, which is indicative of the homogeneous crystal form. Apparently this result is consistent with the crystalline morphology of isotactic polypropylene observed during the transcrystallization, as reported by Varga et al.56 Moreover, we point out that the intensity curves for spherulite and TC basically show four distinct diffraction peaks located at 2θ = 15.0°, 16.9°, 19.3°, and 22.5°, which are attributed to lattice planes (010), (200)/(110), (203) and (015) of α crystal form of PLLA,57 respectively. These data emphasize the impressive fact that the use of heterogeneous fibers or a crystallization accelerator will not affect the way to pack into lamellae for PLLA chains, and thus the crystal form is not encountered with any alteration. SEM observation was further performed to examine the lamellar morphology of TC. Figure 9 provides direct

Figure 10. Typical load−strain curves obtained from single fiber pullout tests of (a) amorphous PLLA by quenching in liquid nitrogen, (b) PLLA3 crystallized at 125 °C, (c) PLLA5 crystallized at 130 °C, and (d) PLLA10 crystallized at 135 °C.

shape of all these curves is similar to that measured by Bannister et al.48 during the pull-out test of single aramid fiber in epoxy matrix, presenting an elastic deformation during the low strain of less than 4%. The maximum debonding forces cannot be used directly for comparison due to the different diameters of embedded ramie fibers (10−40 μm) in the PLLA matrix. To evaluate the interfacial properties, interfacial shear strength (τ) was calculated using the following equation:

Figure 9. Typical SEM micrographs of spherulite and TC formed in the pure PLLA and PLLA3 samples after isothermal crystallization for 30 min at 125 °C. (a) Overview of spherulite and TC around ramie fiber in pure PLLA sample; (b) morphology of spherulite in the bulk matrix of pure PLLA sample; (c) overview of spherulite and TC around ramie fiber in PLLA3 sample; (d) local morphology of TC in PLLA3 sample.

τ= information on lamellar structure of TC and spherulites in the pure PLLA and PLLA3 samples fostered at 125 °C. Figure 9a presents an overview of spherulite and TC around ramie fiber in pure PLLA sample. Obviously the transcrystalline lamellae grow radially whether near the fiber surface or in the impinging front encountered with the adjacent spherulites, like the way that spherulites organize their lamellar growth as depicted in Figure 9b. Similarly, the fan-shaped TC constituting the present columnar morphology indicates no change in the lamellar structure as illustrated in Figure 9c,d. Generally, once nucleated at the fiber surface, TC lamellae radially spread out to the bulk rather than developing perpendicular to the fiber, and hemispheric transcrystalline morphology is formed. Therefore, the SEM observations and 2D-WAXD patterns yield the same results in which the transcrystalline lamellae grow in the same

Fmax πdlm

(1)

where Fmax is the maximum pull-out strength, d represents the fiber diameter determined by optical microscope, and lm indicates the length of the fiber embedded in the PLLA matrix, which is set as 1 mm. The interfacial shear strength for the quenched samples and crystallized samples containing TC textures is presented in Figure 11. The series of amorphous samples that experienced quenching in liquid nitrogen obviously give the lowest strength values, moderately increasing from 5.5 MPa for pure PLLA to 6.4, 6.2, and 5.9 MPa for PLLA3, PLLA5, and PLLA10, respectively. The slight promotion of interfacial shear strength for modified PLLA samples is probably derived from the enhanced interfacial polar interactions between ramie fibers and the introduced PEG chains. The interfacial shear strength H



Article

Figure 12. Schematic representation for the comparison of the crystalline morphology developed in (a) PLLA biocomposite and (b) modified PLLA biocomposite. Note that PEG chains with low molecular weight are miscible with amorphous PLLA.

Figure 11. Interfacial shear strength for ramie fibers partially embedded in pure PLLA, PLLA3, PLLA5, and PLLA10 samples, which are amorphous and isothermally crystallized at 125, 130, and 135 °C, respectively.

the interface.54,55,62 At this stage, the growth of lamellae is spatially hindered by neighboring lamellae, and, finally, the columnar transcrystalline growth proceeds until the growing front impinges with spherulites nucleated in the bulk and encloses the interface.54,55,62 Obviously, the following two prerequisites are critical for the formation of TC: spatial hindrance induced by the established heterogeneous nuclei with sufficiently high density along the interface, and appropriate crystallization conditions that favor the lamellar growth in the direction of fiber radial, but suppress the nucleation in the bulk matrix.53 Many factors such as the matrix type, crystallization conditions, and the fiber material type are considered to affect transcrystalline layers to some extent.42,54,55,63 In our case, judging from the direct observations, it is essentially derived from the fact that ramie fibers can provide relatively adequate nucleating sites like other type of fibers such as polypropylene fiber,64 carbon fiber,39,62,63 carbon nanotube fiber,41 and aramid fiber.52 More importantly, introducing PEG into PLLA absolutely accelerates chain mobility, the crystallization process, and thereby alters the crystalline morphology, as presented in Figure 10b. That is, PLLA obtains preferential ramie fiber-induced nucleation at the fiber surface and rapid transcrystallization with the aid of PEG. Once the heterogeneous nucleating sites are generated, the growth of TC will be substantially assisted by flexible PEG chains, showing independence from the PEG content and crystallization temperature (Figure 5). Of immense significance, too, is the fact that extremely promoted molecular mobility of PLLA attributed to the increased free volume in the presence of short PEG chains, and the disturbance of PEG phase imposes a great challenge on fulfilling the critical size for stable nuclei. It thus permits more valid space for transcrystallization. Therefore we reasonably acquire the understanding that the distinct advantages of using PEG not only come from more valid space for transcrystallization but also stem from the enhanced growth rates of TC. Provided with the prerequisite priorities, TC developed close to the fiber easily wins the competitive growth with the spherulites, and dominates the whole observation area. The rapid generation of thick TC with large size wrapping ramie fibers holds great potential in enhancing their bonding with the biopolymer matrix. Figure 12 also provides insights into the lamellar structure of ramie fiber-induced TC formed in pure PLLA and PLLA/PEG system, presenting fan-shaped morphology and randomly oriented lamellae. Further work

begins a steep increase for the crystallized samples containing TC superstructure. Pure PLLA samples fostered at 125, 130, and 135 °C obtain the interfacial shear strength of 8.3, 9.7, and 8.6 MPa, which rapidly rises to 12.6, 15.4, and 14.5 MPa for the PLLA3 samples, then further climbs to 16.3, 19.5, and 18.3 MPa for the PLLA5 samples, and finally slightly drops down to 14.2, 17.5, and 16.2 MPa for the PLLA10 samples, respectively. It is clear that the PLLA5 samples obtain optimal performances, which lie in the prevailing formation of large-sized TC and simultaneous inhibition of spherulites. For the PLLA10 samples, the addition of 10 wt % PEG reduces the intrinsic strength of the PLLA matrix and pushes down the interfacial properties. Essentially, an appropriate content of the accelerator of chain mobility should be carefully considered for the sake of favorable interfacial adhesion. In general, an accelarator of chain mobility should faciliate the transcrystallization, but not affect the intrinstic strength of matrix too much, which is proved to be mainly dependent on the crystallization temperature and content of PEG. Promisingly, our preliminary attempts shape a green approach to precisely design interfacial crystalline morphology. Recognizing these interesting results, i.e., the generation of prevailing TC in the PLLA/PEG system by comparison to pure PLLA and its significant contribution to enhancing the interfacial bonding, we are stimulated to understand the underlying mechanism associated with the formation of ramie fiber-induced TC in pure and modified PLLA samples. Figure 12 schematically compares the crystalline morphology generated in the pure and PLLA/PEG samples. During crystallization at a relatively low temperature (e.g., 125−135 °C) as shown in Figure 12a, ramie fiber induces heterogeneous nucleation in advance for the subsequent formation of columnar TC at the fiber surface. Nevertheless, poor crystallization kinetics of PLLA itself and severe spatial hindrance from dense neighboring spherulites drag the transcrystallization into a dilemma, ultimately producing limited, thin transcrystalline layers. As previously discussed, during the process of TC formation, a large amount of small crystalline aggregates with relatively imperfect structure develops first from adjacent chains in the interfaces, which can be viewed as primary nuclei to anchor the folded chain crystallites that induce lateral lamellar growth perpendicular to I



■

ACKNOWLEDGMENTS The authors gratefully thank the financial support from the National Natural Science Foundation of China (Grants 51120135002, 51121001, 51203104, and 50925311), the Doctoral Program of the Ministry of Education of China (20130181130012), and the Open Research Fund of State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences (Grant No. 201014). We are in great debt for the kind help with synchrotron 2D X-ray measurements from the beamline BL15U1 and BL16B1 of the Shanghai Synchrotron Radiation Facility (SSRF, Shanghai, China). We appreciate helpful and thoughtful discussions with Prof. Benjamin S. Hsiao from the Department of Chemistry at Stony Brook University.

would fuel theoretical and technological interests in interfacial crystalline structure induced by natural fiber and thus interfacial properties.

■

CONCLUSIONS AND OUTLOOK

Transcrystalline morphology of PLLA induced by ramie fiber was conveniently and precisely tailored by adding an accelerator of chain mobility with short chain length (i.e., PEG) and adequately controlling crystallization temperature. The interfacial morphology and properties were qualitatively and quantitatively studied through direct evidence based on scanning microbeam 2D-WAXD measurements, POM and SEM observations, and single-fiber pull-out test. The utilization of PEG held great effectiveness and power to accelerate the TC growth at ramie fiber during isothermal crystallization. Generally, PLLA was difficult to develop stable nucleation at both high crystallization temperature and high PEG content. In this case, ramie fiber tended to provide effective nucleation sites, and prevailing TC superstructures rather than spherulites were obtained for the PLLA/PEG samples in which the PEG content is inversely proportional to the nucleation density of spherulites. Ultimately, the PLLA/PEG samples containing dominant TC and little spherulite achieved significantly higher interfacial shear strength compared to that of pure PLLA samples. Therefore, it shows high potential to controllably create PLLA biocomposites mainly (or exclusively) comprising transcrystalline structure, permitting a nice chance to tailor mechanical properties by precisely designing TC texture in PLLA biocomposites. Furthermore, it is instructive to realize the synergetic effect of ramie fiber and PEG on the development of TC with respect to guiding the practical processing, where ramie fibers elevate the nucleation rate in the upper temperature window and PEG profoundly improves the molecular mobility in the lower temperature range. As a result, it is reasonable that the combination of an effective nucleating agent and a crystallization modifier will significantly broaden the crystallization window. This distinguished feature will be helpful when we are faced with controlling the crystalline morphology in the industrial fabrication of semicrystalline polymer composites, especially for polymers with low crystallization ability such as PLLA. Of immense significance, too, are green features that have been inherently embedded in the present approach, but previous reports have thus far been unable to realize the enhanced interfacial properties by modifying the surfaces of natural fibers, such as physical and chemical treatment covering a wide range from microwaving, mercerization, surface-grafting, coupling, to alkalization, which undesirably means extreme time consumption, energy wasting, and the consumption of tons of noxious and toxic reagents.

■

Article

■

REFERENCES

(1) Zhong, Y.; Fang, H.; Zhang, Y.; Wang, Z.; Yang, J.; Wang, Z. Rheologically Determined Critical Shear Rates for Shear-Induced Nucleation Rate Enhancements of Poly(lactic acid). ACS Sustainable Chem. Eng. 2013, 1, 663−672. (2) de Aguiar, H. B.; de Beer, A. G. F.; Roke, S. The Presence of Ultralow Densities of Nanocrystallites in Amorphous Poly(lactic acid) Microspheres. J. Phys. Chem. B 2013, 117, 8906−8910. (3) Zhou, S.; Peng, H.; Yu, X.; Zheng, X.; Cui, W.; Zhang, Z.; Li, X.; Wang, J.; Weng, J.; Jia, W.; Li, F. Preparation and Characterization of a Novel Electrospun Spider Silk Fibroin/Poly(D,L-lactide) Composite Fiber. J. Phys. Chem. B 2008, 112, 11209−11216. (4) Drumright, R. E.; Gruber, P. R.; Henton, D. E. Polylactic Acid Technology. Adv. Mater. 2000, 12, 1841−1846. (5) Daiguji, H.; Takada, S.; Cornejo, J. J. M.; Takemura, F. Fabrication of Hollow Poly(lactic acid) Microcapsules from Microbubble Templates. J. Phys. Chem. B 2009, 113, 15002−15009. (6) Delpouve, N.; Stoclet, G.; Saiter, A.; Dargent, E.; Marais, S. Water Barrier Properties in Biaxially Drawn Poly(lactic acid) Films. J. Phys. Chem. B 2012, 116, 4615−4625. (7) Zhao, Z.-J.; Wang, Q.; Zhang, L.; Liu, Y.-C. A Different Diffusion Mechanism for Drug Molecules in Amorphous Polymers. J. Phys. Chem. B 2007, 111, 4411−4416. (8) Molino Cornejo, J. J.; Daiguji, H.; Takemura, F. Factors Affecting the Size and Uniformity of Hollow Poly(lactic acid) Microcapsules Fabricated from Microbubble Templates. J. Phys. Chem. B 2011, 115, 13828−13834. (9) Kasimova, A. O.; Pavan, G. M.; Danani, A.; Mondon, K.; Cristiani, A.; Scapozza, L.; Gurny, R.; Möller, M. Validation of a Novel Molecular Dynamics Simulation Approach for Lipophilic Drug Incorporation into Polymer Micelles. J. Phys. Chem. B 2012, 116, 4338−4345. (10) Dorgan, J. R., Braun, B., Wegner, J. R., Knauss, D. M. Poly(lactic acids): A brief review. In Degradable Polymers and Materials. Principles and Practice; Khemani, K., Scholz, C., Eds.; American Chemical Society: Washington, D.C., 2006; pp 102−125. (11) Haspel, N.; Laurent, A. D.; Zanuy, D.; Nussinov, R.; Alemán, C.; Puiggalí, J.; Revilla-López, G. Conformational Exploration of Two Peptides and Their Hybrid Polymer Conjugates: Potentialities as SelfAggregating Materials. J. Phys. Chem. B 2012, 116, 13941−13952. (12) Gao, M.; Ren, Z.; Yan, S.; Sun, J.; Chen, X. An Optical Microscopy Study on the Phase Structure of Poly(L-lactide acid)/ Poly(propylene carbonate) Blends. J. Phys. Chem. B 2012, 116, 9832− 9837. (13) Bao, R.-Y.; Yang, W.; Jiang, W.-R.; Liu, Z.-Y.; Xie, B.-H.; Yang, M.-B. Polymorphism of Racemic Poly(L-lactide)/Poly(D-lactide) Blend: Effect of Melt and Cold Crystallization. J. Phys. Chem. B 2013, 117, 3667−3674. (14) Aou, K.; Hsu, S. L.; Kleiner, L. W.; Tang, F.-W. Roles of Conformational and Configurational Defects on the Physical Aging of Amorphous Poly(lactic acid). J. Phys. Chem. B 2007, 111, 12322− 12327.

AUTHOR INFORMATION

Corresponding Authors

*Ph +86-28-8540-0211; Fax +86-28-8540-6866; e-mail: ganji. [email protected] (G.-J.Z.). *Ph +86-28-8540-6866; Fax +86-28-8540-6866; e-mail: zmli@ scu.edu.cn (Z.-M.L.). Notes

The authors declare no competing financial interest. J



Article

(15) Lu, J.; Qiu, Z.; Yang, W. Fully Biodegradable Blends of Poly(Llactide) and Poly(ethylene succinate): Miscibility, Crystallization, and Mechanical Properties. Polymer 2007, 48, 4196−4204. (16) Xu, H.; Xie, L.; Chen, Y.-H.; Huang, H.-D.; Xu, J.-Z.; Zhong, G.J.; Hsiao, B. S.; Li, Z.-M. Strong Shear Flow-Driven Simultaneous Formation of Classic Shish-Kebab, Hybrid Shish-Kebab and Transcrystallinity in Poly(lactic acid)/Natural Fiber Biocomposites. ACS Sustainable Chem. Eng. 2013, DOI: 10.1021/sc4003032. (17) Ojijo, V.; Sinha Ray, S.; Sadiku, R. Toughening of Biodegradable Polylactide/Poly(butylene succinate-co-adipate) Blends via in Situ Reactive Compatibilization. ACS Appl. Mater. Interfaces 2013, 5, 4266−4276. (18) Ojijo, V.; Sinha Ray, S.; Sadiku, R. Effect of Nanoclay Loading on the Thermal and Mechanical Properties of Biodegradable Polylactide/Poly [(butylene succinate)-co-adipate] Blend Composites. ACS Appl. Mater. Interfaces 2012, 4, 2395−2405. (19) Ljungberg, N.; Wesslén, B. Tributyl Citrate Oligomers as Plasticizers for Poly(lactic acid): Thermo-Mechanical Film Properties and Aging. Polymer 2003, 44, 7679−7688. (20) Kulinski, Z.; Piorkowska, E. Crystallization, Structure and Properties of Plasticized Poly(L-lactide). Polymer 2005, 46, 10290− 10300. (21) Kulinski, Z.; Piorkowska, E.; Gadzinowska, K.; Stasiak, M. Plasticization of Poly(L-lactide) with Poly(propylene glycol). Biomacromolecules 2006, 7, 2128−2135. (22) Hu, Y.; Hu, Y. S.; Topolkaraev, V.; Hiltner, A.; Baer, E. Aging of Poly(lactide)/Poly(ethylene glycol) Blends. Part 2. Poly(lactide) with High Stereoregularity. Polymer 2003, 44, 5711−5720. (23) Lai, W.-C.; Liau, W.-B.; Lin, T.-T. The Effect of End Groups of PEG on The Crystallization Behaviors of Binary Crystalline Polymer Blends PEG/PLLA. Polymer 2004, 45, 3073−3080. (24) Pillin, I.; Montrelay, N.; Grohens, Y. Thermo-Mechanical Characterization of Plasticized PLA: Is the Miscibility the Only Significant Factor? Polymer 2006, 47, 4676−4682. (25) Ljungberg, N.; Wesslén, B. Preparation and Properties of Plasticized Poly(lactic acid) Films. Biomacromolecules 2005, 6, 1789− 1796. (26) Kumar, S.; Hofmann, M.; Steinmann, B.; Foster, E. J.; Weder, C. Reinforcement of Stereolithographic Resins for Rapid Prototyping with Cellulose Nanocrystals. ACS Appl. Mater. Interfaces 2012, 4, 5399−5407. (27) Summerscales, J.; Dissanayake, N. P. J.; Virk, A. S.; Hall, W. A Review of Bast Fibres and Their Composites. Part 1 − Fibres As Reinforcements. Compos. Part A: Appl. Sci. 2010, 41, 1329−1335. (28) Chang, C.-W.; Wang, M.-J. Preparation of Microfibrillated Cellulose Composites for Sustained Release of H2O2 or O2 for Biomedical Applications. ACS Sustainable Chem. Eng. 2013, 1, 1129− 1134. (29) Nagarajan, V.; Mohanty, A. K.; Misra, M. Sustainable Green Composites: Value Addition to Agricultural Residues and Perennial Grasses. ACS Sustainable Chem. Eng. 2013, 1, 325−333. (30) Bledzki, A. K.; Jaszkiewicz, A. Mechanical Performance of Biocomposites Based on PLA and PHBV Reinforced with Natural Fibres − A Comparative Study to PP. Compos. Sci. Technol. 2010, 70, 1687−1696. (31) Nagarajan, V.; Mohanty, A. K.; Misra, M. Sustainable Green Composites: Value Addition to Agricultural Residues and Perennial Grasses. ACS Sustainable Chem. Eng. 2013, 1, 325−333. (32) Chung, Y.-L.; Olsson, J. V.; Li, R. J.; Frank, C. W.; Waymouth, R. M.; Billington, S. L.; Sattely, E. S. A Renewable Lignin−Lactide Copolymer and Application in Biobased Composites. ACS Sustainable Chem. Eng. 2013, DOI: 10.1021/sc4000835. (33) Wang, T.; Drzal, L. T. Cellulose-Nanofiber-Reinforced Poly(lactic acid) Composites Prepared by a Water-Based Approach. ACS Appl. Mater. Interfaces 2012, 4, 5079−5085. (34) Iman, M.; Bania, K. K.; Maji, T. K. Green Jute-Based CrossLinked Soy Flour Nanocomposites Reinforced with Cellulose Whiskers and Nanoclay. Ind. Eng. Chem. Res. 2013, 52, 6969−6983.

(35) Xu, H.; Liu, C.-Y.; Chen, C.; Hsiao, B. S.; Zhong, G.-J.; Li, Z.-M. Easy Alignment and Effective Nucleation Activity of Ramie Fibers in Injection-Molded Poly(lactic acid) Biocomposites. Biopolymers 2012, 97, 825−839. (36) Quan, Y.; Li, H.; Yan, S. Comparison Study on the Heterogeneous Nucleation of Isotactic Polypropylene by Its Own Fiber and α Nucleating Agents. Ind. Eng. Chem. Res. 2013, 52, 4772− 4778. (37) Feldman, A. Y.; Wachtel, E.; Vaughan, G. B. M.; Weinberg, A.; Marom, G. The Brill Transition in Transcrystalline Nylon-66. Macromolecules 2006, 39, 4455−4459. (38) Cho, K.; Kim, D.; Yoon, S. Effect of Substrate Surface Energy on Transcrystalline Growth and Its Effect on Interfacial Adhesion of Semicrystalline Polymers. Macromolecules 2003, 36, 7652−7660. (39) Zhang, M.; Xu, J.; Zhang, Z.; Zeng, H.; Xiong, X. Effect of Transcrystallinity on Tensile Behaviour of Discontinuous Carbon Fibre Reinforced Semicrystalline Thermoplastic Composites. Polymer 1996, 37, 5151−5158. (40) Folkes, M. J.; Hardwick, S. T. Direct Study of the Structure and Properties of Transcrystalline Layers. J. Mater. Sci. Lett. 1987, 6, 656− 658. (41) Zhang, S.; Minus, M. L.; Zhu, L.; Wong, C.-P.; Kumar, S. Polymer Transcrystallinity Induced by Carbon Nanotubes. Polymer 2008, 49, 1356−1364. (42) Klein, N.; Marom, G. Microstructure of Nylon 66 Transcrystalline Layers in Carbon and Aramid Fibre Reinforced Composites. Polymer 1996, 37, 5493−5498. (43) Campbell, D.; Qayyum, M. M. Enhanced Fracture Strain of Polypropylene by Incorporation of Thermoplastic Fibres. J. Mater. Sci. 1977, 12, 2427−2434. (44) Felix, J. M.; Gatenholm, P. Effect of Transcrystalline Morphology in Interfacial Adhesion in Cellulose-PP Composites. J. Mater. Sci. 1994, 29, 3043−3049. (45) Wang, Y.; Tong, B.; Hou, S.; Li, M.; Shen, C. Transcrystallization Behavior at the Poly(lactic acid)/Sisal Fibre Biocomposite Interface. Composites, Part A 2011, 42, 66−74. (46) De Marco, R.; Tolomelli, A.; Greco, A.; Gentilucci, L. Controlled Solid Phase Peptide Bond Formation Using NCarboxyanhydrides and PEG Resins in Water. ACS Sustainable Chem. Eng. 2013, 1, 566−569. (47) Chen, Y.-H.; Zhong, G.-J.; Wang, Y.; Li, Z.-M.; Li, L. Unusual Tuning of Mechanical Properties of Isotactic Polypropylene Using Counteraction of Shear Flow and β-Nucleating Agent on β-Form Nucleation. Macromolecules 2009, 42, 4343−4348. (48) Bannister, D.; Andrews, M.; Cervenka, A.; Young, R. Analysis of the Single-Fibre Pull-Out Test by Means of Raman Spectroscopy: Part II. Micromechanics of Deformation for an Aramid/Epoxy System. Compos. Sci. Technol. 1995, 53, 411−421. (49) Xu, H.; Zhong, G.-J.; Fu, Q.; Lei, J.; Jiang, W.; Hsiao, B. S.; Li, Z.-M. Formation of Shish-Kebabs in Injection-Molded Poly(L-lactic acid) by Application of an Intense Flow Field. ACS Appl. Mater. Interfaces 2012, 4, 6774−6784. (50) Hikosaka, M. Unified Theory of Nucleation of Folded-Chain Crystals and Extended-Chain Crystals of Linear-Chain Polymers. Polymer 1987, 28, 1257−1264. (51) Moon, C.-K. The Effect of Interfacial Microstructure on the Interfacial Strength of Glass Fiber/Polypropylene Resin Composites. J. Appl. Polym. Sci. 1994, 54, 73−82. (52) Assouline, E.; Wachtel, E.; Grigull, S.; Lustiger, A.; Wagner, H. D.; Marom, G. Lamellar Orientation in Transcrystalline γ Isotactic Polypropylene Nucleated on Aramid Fibers. Macromolecules 2002, 35, 403−409. (53) Quan, H.; Li, Z.-M.; Yang, M.-B.; Huang, R. Review on Transcrystallinity in Semi-crystalline Polymer Composites. Compos. Sci. Technol. 2005, 65, 999−1021. (54) Li, H.; Yan, S. Surface-Induced Polymer Crystallization and the Resultant Structures and Morphologies. Macromolecules 2011, 44, 417−428. K



Article

(55) Ninomiya, N.; Kato, K.; Fujimori, A.; Masuko, T. Transcrystalline Structures of Poly(L-lactide). Polymer 2007, 48, 4874−4882. (56) Varga, J.; Karger-Kocsis, J. Direct Evidence of Row-Nucleated Cylindritic Crystallization in Glass Fiber-Reinforced Polypropylene Composites. Polym. Bull. 1993, 30, 105−110. (57) Zhang, J.; Duan, Y.; Sato, H.; Tsuji, H.; Noda, I.; Yan, S.; Ozaki, Y. Crystal Modifications and Thermal Behavior of Poly(L-lactic acid) Revealed by Infrared Spectroscopy. Macromolecules 2005, 38, 8012− 8021. (58) Romero, I. S.; Schurr, M. L.; Lally, J. V.; Kotlik, M. Z.; Murphy, A. R. Enhancing the Interface in Silk−Polypyrrole Composites through Chemical Modification of Silk Fibroin. ACS Appl. Mater. Interfaces 2013, 5, 553−564. (59) Magniez, K.; Voda, A. S.; Kafi, A. A.; Fichini, A.; Guo, Q.; Fox, B. L. Overcoming Interfacial Affinity Issues in Natural Fiber Reinforced Polylactide Biocomposites by Surface Adsorption of Amphiphilic Block Copolymers. ACS Appl. Mater. Interfaces 2012, 5, 276−283. (60) Nordgren, N.; Carlsson, L.; Blomberg, H.; Carlmark, A.; Malmström, E.; Rutland, M. W. Nanobiocomposite Adhesion: Role of Graft Length and Temperature in a Hybrid Biomimetic Approach. Biomacromolecules 2013, 14, 1003−1009. (61) Quero, F.; Nogi, M.; Yano, H.; Abdulsalami, K.; Holmes, S. M.; Sakakini, B. H.; Eichhorn, S. J. Optimization of the Mechanical Performance of Bacterial Cellulose/Poly(L-lactic) Acid Composites. ACS Appl. Mater. Interfaces 2009, 2, 321−330. (62) Varga, J.; Karger-Kocsis, J. Interfacial Morphologies in Carbon Fiber-Reinforced Polypropylene Microcomposites. Polymer 1995, 36, 4877−4881. (63) Thomason, J. L.; Van Rooyen, A. A. Transcrystallized Interphase in Thermoplastic Composites. J. Mater. Sci. 1992, 27, 889−896. (64) Li, H.; Zhang, X.; Duan, Y.; Wang, D.; Li, L.; Yan, S. Influence of Crystallization Temperature on the Morphologies of Isotactic Polypropylene Single-Polymer Composite. Polymer 2004, 45, 8059− 8065.

L


Natural

Recommend Documents