Polypropylene

Research Article pubs.acs.org/journal/ascecg

Cite This: ACS Sustainable Chem. Eng. 2017, 5, 10413-10420

Enhanced Interfacial Strength of Natural Fiber/Polypropylene Composite with Mechanical-Interlocking Interface Shaohong Shi, Changhua Yang, and Min Nie* State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of Sichuan University, No. 24 South Section 1, Yihuan Road, Chengdu 610065, China

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S Supporting Information *

ABSTRACT: The interface, as the weakest point of a polymer composite, determines the comprehensive performance, especially in a polarity/nonpolarity system featuring poor interfacial adhesion. Here, we report an interfacial manipulation strategy to trigger the preferential adsorption of amide-based self-assembling compounds (NAs) from polypropylene (PP) melts onto the surface of natural fiber (NF) as a result of hydrogen bonding and then promote the epitaxial growth, into root-like NF fiber with interfacial interlocking effects. The unique interface constructed by the grown NA fibers rendered the PP/NF composite with strong interfacial adhesion. The substantial increases of 64.4%, 77.9%, and 94.4% in interfacial shear strength, interfacial friction, and debonding energy are achieved, respectively, in comparison to conventional NF/PP composite. Finally, the working principle of the laterally grown NA fibers on the interfacial enhancement was established based on the fracture morphology after the microbond test. This study can effectively solve the interfacial problems of a polymer composite featuring limited interfacial adhesion, via simple one-step physical blending, without any preliminary surface treatment or “soft” compatibilizers. KEYWORDS: Polypropylene, Natural fiber, Mechanical interlocking, Interfacial adhesion, self-assembly

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INTRODUCTION Polymer composites possess an integrated combination of structures and properties from the polymer matrix and functional fillers and thus hold the potential as high-strength and functional materials.1−4 In this field, an important topic is fiber-reinforced polymer composite. Ascribing to the renewability, sustainability, biodegradability, abundant availability, and outstanding mechanical properties, natural fiber (NF) is considered as an ideal green reinforcing biofiller for polymer, evoking considerable interest in preparing polymer/NF composites.5−8 The reinforcing fibers come to play via effective load-transfer from the polymer matrix to the fiber via the interface.9 Excellent interfacial adhesion plays a crucial role in guaranteeing the reinforcing efficiency of the fiber and thereby the bulk mechanical performances of the resulting composites. Therefore, successful utilization of NFs for the reinforcing polymer not only is dependent on their intrinsic structures and mechanical properties but also is equally contributed by the fiber−matrix interface.10−12 However, NFs often exhibit inherent incompatibility with most polymers, especially for nonpolar polymer including polypropylene (PP), leading to weak interfacial adhesion and limited reinforcement efficiency.13−15 As a result, the NF/PP composites prepared via conventional process have inferior mechanical properties. Polymer/fiber interfacial design and optimization are the long-standing critical issues for developing high-strength © 2017 American Chemical Society

composites. One common protocol employed for enhancing the interfacial adhesion is to minimize the unfavorable polymerfiber interfacial energy, such as modifying the filler surface via chemical grafting method and incorporating some compatibilizers with dual functionalities.16−19 Although these strategies have been applied successfully in NF/polymer composite, some drawbacks are still not solved well, such as the complicated synthesis process accompanied by vast consumption of toxic reagents, decreased bulk strength of the composites induced by “soft” compatibilizer featuring low-molecular weight, and so on.20−22 Besides lowering interfacial energy, enhancing interfacial friction to restrict mutual sliding of two phases is another modifying way for improving the interfacial adhesion.23−25 This has been well exemplified by some inorganic fillers featuring unique geometry,26−28 such as tetrapodal-shaped zinc oxide microparticle, bamboo-like silica fibers, shish-kebab carbon nanotube, etc. In the cases, the fillers interlock mechanically with the polymer matrix, generating high interfacial friction to facilitate the load transfer efficiency via interface.28,29 Recently, upon the inspiration that it is difficult to pull out a tree root whose branches tightly embed into their surrounding soil, rootlike fibers are produced by some inorganic fibers growing on Received: July 20, 2017 Revised: September 7, 2017 Published: September 21, 2017 10413

DOI: 10.1021/acssuschemeng.7b02448 ACS Sustainable Chem. Eng. 2017, 5, 10413−10420

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the surface of the reinforcing fibers to mechanically interlock with a polymer matrix.30−32 Accordingly, effective load transfer and good interfacial adhesion are achieved to perfect the resulting composites. Tamrakar grafted carbon nanotubes (CNTs) onto micrometer-sized glass fiber (GF) via electrophoretic deposition and the interfacial shear strength between the CNTs-g-GF, and epoxy resin increased by 58% relative to the pristine GF.33 Patterson adopted an ion-exchange technique to allow zinc oxide nanowires to be directly deposited on the surface of aramid fibers to improve the mechanical properties of the fiber-reinforced epoxy composite with 33.3% increase of interfacial strength.34 Now, the root-like fibers have been applied successfully to many thermoset polymers but fail when applied to thermoplastic polymers because the branched fibers can be peeled away from the primary fibers during meltcompounding of the fibers and the corresponding polymers, where large shear stresses are generated during twin-screw extrusion.35−38 Very recently, our groups proposed a facile and effective way to directly construct root-like fiber in polymer melts, involving the interfacial diffusion of some amide-based self-assembling compounds (NAs) driven by the polarity differences between the components and the subsequent epitaxial crystallization on the surface of the reinforcing filler. As a result, the grown fibers at the interface functioned as interlocks to achieve the mechanical interlocking of the filler with the polymer matrix.39,40 However, the amide-based selfassembling compounds are polar molecules so that the interfacial diffusion only appears when the surface nature of the fillers is transformed into the similar polarity via additional chemical modification; moreover, the strengthening effects of the interfacially grown fibers on the interfacial adhesion between the fillers and polymer matrix are still far from being understood. Compared to the other reinforcing fillers, numerous hydroxyl groups are loaded at the surface of NF, presenting a large potential to form hydrogen bonding with the amide-based molecules.41 It can be envisioned that NFs can act as an ideal acceptor to anchor and immobilize the amide-based selfassembling compounds, achieving the directed diffusion and in situ construction of root-like NFs, with no need for additional chemical pretreatment on the surface. Accordingly, there are some issues that need to be addressed: (1) How does hydrogen-bonding form between the amide-based compounds and NF? (2) Is the interaction large enough to drive the directed diffusion of the amide-based compounds? (3) How does the root-like fiber debond from polymer matrix? The answer to the above questions should help us establish an advanced green avenue for preparing NF-reinforced polymer composites with excellent interfacial natural and mechanical properties. In this end, we selected a well-established amidebased self-assembling compound (NA) with a good dissolution at elevated temperature and then self-assembly in PP melts upon cooling,42 which is a prerequisite to prepare root-like fiber via the interfacial diffusion and growth. First, the underlying interaction between NAs and NFs was evaluated by twodimensional correlation FTIR spectra, and then the growth process of the root-like NF in the PP melts was investigated. Finally, a single fiber pull-out test was carried out to capture fully the microdebonding mechanism of the root-like NF from the polymer matrix. This study sufficiently demonstrated the simple preparation of root-like NF and the potential of improving interfacial adhesion in NF-reinforced PP composites by regulating interfacial topology.

Research Article

EXPERIMENTAL SECTION

Materials. Isotactic polypropylene (trade name: T30S) with melt flow rate of 2.6 g/10 min was provided by Dushanzi Petroleum Chemical Incorporation (Xinjiang, China). Sisal fiber, a kind of widely used natural fiber (NF), was provided by Raffia Paper Products Company (Zhejiang, China). Rare earth compound (trade name: WBG-II) was purchased from Guangdong Weilinna Incorporation (Guangdong, China). It is a kind of dimetal complex of lanthanum and calcium with dicarboxylic acid and amide-type ligands. The molecular formula is CaxLa1−x(LIG1)m(LIG2)n, where x and 1 − x are the proportion of Ca2+ and La3+ ions in the complex while LIG1 and LIG2 represent dicarboxylic acid and amide-type ligands with coordination numbers of m and n, respectively. Sample Preparation. Natural fibers were dried in vacuum at 40 °C overnight before use to remove the residual waters. NA powders and PP pellets were premixed and extruded by a twin screw extruder with length to diameter (L/D) ratio of 32. The screw speed was 50 r/min and processing temperature was 180 °C. The NAs concentration was 0.3%. Next, NA-containing PP samples were hotpressed into ∼10 μm thick films for the following characterizations. Characterization. Fourier Transform Infrared Spectroscopy (FTIR) and Two-Dimensional Correlation Spectroscopy (2DCS) Analysis. NA powders were mixed with NFs at the weight ratio of 1:1, put into KBr powders, and pressed into a disk for the FT-IR investigation. The FTIR spectra of NA, NF, and the NA/NF mixture were collected over the wavenumber range of 4000−500 cm−1 by a Nicolet 20SXB spectrometer (Thermo Fisher Scientific Inc., U.S.A.). In addition, in order to clearly reveal the interaction between NF and NA, the temperature-dependent FTIR spectra between 30 and 200 °C were also conducted at a heating rate of 5 °C/min. Then, twodimensional correlation spectroscopy (2DCS) analysis was performed in the region of 3000−3600 cm−1 by using the 2DCS 6.1 software developed by Tao Zhou at Sichuan University. In the 2D spectra, red color regions are defined as positive correlation, whereas blue color regions are regarded as negative correlation.43 Polarized Light Microscope (PLM) Observation. For convenient observation on the interfacial structure, single NF was directly embedded between two NA-containing PP films. The assembly was observed under polarized light microscope (Leica DM2500P) equipped with hot stage (Linkam THMS600, Linkam Scientific Instruments Ltd., U.K.). The sample was heated to 230 °C and kept for 5 min and then cooled down to 140 °C at a rate of 10 °C/min. The interface structures at the NF surface and the subsequent crystallization process of PP were recorded. Single Fiber Pull-Out Test. The single fiber pull-out test was adopted to evaluate the interfacial properties of the NF/PP composite. The samples after PLM investigation were measured by an Instron universal test instrument (Model 5576, Instron Instruments, U.S.A.) with a gauge length of 10 mm and crosshead speed of 2 mm/min. The imbedded length in the PP matrix is ∼0.5 mm. At least 5 replicates were performed for each measurement, and the average value was calculated. Scanning Electron Microscope (SEM) Observation. A total of ∼0.3 g of rare earth compound (NA) powders were added into 100 mL of dimethyl sulfoxide at 80 °C and mechanically stirred for 6 h. Then, the NFs were put into the NA solution for 30 min, washed with deionized water, and dried in vacuum at 40 °C. The surface morphology was investigated by an FEI Inspect F-SEM instrument with an acceleration voltage of 20 kV. Additionally, the NF/NAs-containing PP and NF/PP samples after PLM investigation were etched in permanganate etchant at 40 °C for 0.5 h and then observed after being gold-sputtered. To reveal the interfacial deboning mechanism, the polymer matrix and NF surface were also observed after the single fiber pull-out test.

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RESULTS AND DISCUSSION First, the immobilization of NA on the NF was visualized via SEM examination, where NF was dip-coated by NA followed by solvent washing. As shown in Figure 1, compared to the smooth surface of pristine NF, the dip-coating NF, even upon 10414

DOI: 10.1021/acssuschemeng.7b02448 ACS Sustainable Chem. Eng. 2017, 5, 10413−10420

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cm−1. However, the shifted peak was located between the characterized peaks of hydroxyl groups of NF and free NH groups from NA. The broad strong hydroxyl peak may overlap partially with the neighboring subpeak, giving rise to additional peak shift. Therefore, based on the common FTIR results, it is not concluded convincingly that there exists hydrogen bonding between NF and NA. The generalized two-dimensional correlation FTIR spectroscopy (2DCOS) is one powerful technique to distinguish different kinds of overlapped absorption peaks from various groups based on external perturbation-induced signal fluctuations.46,47 Environment temperature can exert a remarkable effect on the association or disassociation of the hydrogen bond, i.e., at high temperature, the hydrogen bonds can be weakened or disassociated.48,49 Accordingly, in this study, temperature was chosen as the external perturbation resource, and the temperature-dependent FTIR spectra of the NA/NF mixture were investigated. Figure 2b presents the changes of the characteristic peaks in the wavenumber range of 3600− 3000 cm−1. One can observe that, with the increasing temperature, the absorption peaks at 3415 and 3317 cm−1 took a shift to higher frequency and the intensity became weak. Further, the peak intensity at 3415 cm−1 was plotted against the temperature in Figure 2c. The intensity first decreased slowly and then dropped quickly. The obvious transition implies that the peak shift at 3415 cm−1 is not resulted from the overlapping of the original hydroxyl and free NH peaks but should be attributed to the hydrogen bonds.50 Obviously, the destruction of the hydrogen bonds during the heating process is responsible for the intensity transition. This also can be verified by the corresponding two-dimensional correlation analysis, as shown in Figure 2d−e. In the synchronous correlation spectra, two auto peaks featuring positive correlation appeared at 3317 and 3415 cm−1, indicating the two hydrogen bonds existed in the NA/NF mixture and the coincidental changes happened; in the asynchronous spectra, the emergence of the cross peaks further revealed that the groups assigned to 3317 and 3415 cm−1 were

Figure 1. SEM micrographs of pristine NF before (a, a1) and after (b, b1, b2) dip-coating treatment with NAs.

the solvent washing, was still wrapped tightly by numerous submicrosized NA granules. It is supposed to result from the underlying interaction between NA and NF, which gets NAs firmly attached to the surface of NF. To verify this assumption, the common FTIR investigations of NA, NF, and NA/NF mixture were conducted first. As shown in Figure 2a, rich hydroxyl groups in NF excited a distinct, broad absorption peak in the range of 3600−3000 cm−1. In the spectrum of NA, the sharp absorption peak at 3317 cm−1 was accompanied by a shoulder one at 3458 cm−1. It is evident that, for amide-based compound, the former is assigned to self-association NH groups while the latter represents free ones.44,45 For NA/NF mixture, the absorption peak corresponding to self-association NH remained unchanged while the shoulder shifted to 3415

Figure 2. FTIR spectra of NA, NF, and the NA/NF mixture at room temperature (a), the temperature-dependent FTIR spectra of the NA/NF mixture in the range of 3600−3000 cm−1 (b), the intensity at 3415 cm−1 as a function of temperature in the range of 30 to 200 °C (c), and the corresponding two-dimensional correlation FTIR spectroscopy of the mixture in the region of 3600−3000 cm−1 (d, e). 10415

DOI: 10.1021/acssuschemeng.7b02448 ACS Sustainable Chem. Eng. 2017, 5, 10413−10420

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Figure 3. PLM of the interfacial evolution of NF/NAs-containing PP upon cooling from 230 °C (a−d): a, 230 °C; b, 180 °C; c, 172 °C; d, 162 °C; isothermal crystallization for 1 min at 140 °C, transcrystalline (TC) is marked by the arrow (e); the sizes of the NA fiber as a function of temperature (f).

Figure 4. SEM micrographs of interfacial morphology of NF/NA-containing PP.

the fibrous NAs alignment was pointed into the PP matrix from the surface of the NF. This is consistent with the PLM in Figure 3. In addition, one can observe that PP crystalline grew around the NA fibers. Ascribed to the existence of the transcrystalline, there is strong interfacial adhesion between polymer and the laterally grown NA fiber,55 which can strengthen the interlocking effect of the root-like NF with the PP matrix. It is well-documented that the effective load transfer from the polymer matrix across the interphase is crucial to bring reinforcing efficiency of a given filler into full play. The mechanical-interlocking interface can integrate the polymer matrix with the filler to improve the interfacial adhesion and thus facilitate the load transfer, which is a favorable manner for developing high-performance polymer composites with poor interfacial adhesion.37,56 However, the response of the unique interface constructed by the interfacially grown NA fibers to external stress remains poorly understood as yet. Here, we adopted a single fiber pull-out test to evaluate the interfacial structure/property relationship, and the corresponding stress− strain curves were plotted in Figure 5. At a given strain, the stress required for a single NF/PP composite featuring the interfacially grown NA fibers always was larger than that of the conventional NF/PP composite, suggesting better interfacial adhesion. When the load was applied to the NF, the pull-out was involved in the interfacial debonding and the friction sliding processes.33,57 The stress−strain curve can be divided into two distinguishable regions: Between points A and B, with

restricted mutually and exhibited a negative correlation. According to Noda’s rule,46,51 it can be deduced that there are two kinds of hydrogen bonds: the low wavenumbers should be assigned to the self-association NH groups while the high one originates from the hydrogen bonding between hydroxyl groups from NF and free NH groups from NA. When NF is introduced into NA-containing PP melts, NAs can preferentially diffuse onto the NF surface driven by the NA−NF hydrogen bonding. Upon cooling, the absorbed NAs can be transformed into the fiber via epitaxial crystallization, forming the root-like NF fiber. As shown in Figure 3a−d, with the decreasing temperature, more NAs recrystallized from the PP melt so the laterally grown NA fiber became bigger and denser. The size of the NA fiber against the temperatures was plotted in Figure 3f. The sigmoid curve indicates that the growth of NA fiber principally proceeded at the narrow temperature range of 182−162 °C and ceased at the lower temperature. This can be attributed to quick recrystallization of polar NAs in the nonpolar PP melts.52,53 In contrast to the weak nucleating efficiency of the pristine NF (see PLM photos of PP/NF composite in Supporting Information), the NA fibers grown at the interface hosted lots of heterogeneous nuclei for PP to enable the epitaxial crystallization at the surface,54 forming the closely wrapping transcrystalline layers. The direct evidence of the transcrystalline layers grown at the surface of NA fibers is visible in Figure 4. Due to poor resistance of the organic molecule NA to chemical etchant relative to long-chain PP, the holes in the SEM images represented the NAs, implying 10416

DOI: 10.1021/acssuschemeng.7b02448 ACS Sustainable Chem. Eng. 2017, 5, 10413−10420

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Figure 5. Typical load−strain curves of NF/NA-containing PP and conventional NF/PP composite during single fiber pull-out test (a1−a3); (b) the calculated interfacial properties including interfacial shear strength (τ), debonding energy (Ecr), and friction stress (τfr).

Figure 6. SEM micrographs of interfacial morphology of the pristine fibers (a), the samples of conventional NF/PP, fiber surface (b, b1) and matrix (c, c1), and NF/NA-containing PP, fiber surface (d, d1) and matrix (e, f, f1), after the single fiber pull-out test.

verified by the debonding energy estimated from the unit area between the points A and B, where more debonding energies were required for the former sample. Beyond point B, a sharp drop of the stress from Fmax to Ffr was observed due to the complete debonding. The remaining stress is attributed by the interfacial friction between the fiber and the matrix. In this case, interfacial friction dictated the displacement of NF during the single fiber pull-out test due to the preventing effects on the NF pull-out. The friction stress (τfr) was calculated based on the remaining stress (Ffr) at point C.57

the increasing displacement, the stress increased and reached the maximum value at the point B; in this region, the crack was initiated at the fiber end and propagated along the fiber surface, finally leading to the complete debonding. The interfacial debonding can be reflected by the interfacial shear strength (τ) and was calculated according to the following equation:58 τ=

Fmax df · π · l f

(1)

where Fmax is the maximum load value, df is the NF diameter, and lf is the length of NF embedded in the matrix. As shown in Figure 5b, τ of NF/PP composite with interfacial interlocking structure was 21.8 MPa, with a noticeable increase of 64.4% compared to that of the conventional NF/PP composite. The interfacially grown NA fibers can effectively enhance the interfacial adhesion between the NF and PP matrix. This is also

τfr =

Ffr df ·π ·lfr

(2)

The interfacial friction increased from 8.4 MPa of conventional NF/PP composite to 15.0 MPa. Many researches have confirmed that the frictional force can promote interfacial 10417

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entrance of the fiber into the matrix first causes the fragments of the NA fibers at the embedded end and initiates the crack. Afterward, the crack propagates along the interface, accompanied by the destruction of the PP matrix with energy dissipation. When the NA fibers are completely destroyed, the fragmented NA granules provide extra friction to persistently resist the debonding process of NF. Accordingly the interfacially grown NA fibers enable the interfacial enhancements. Compared to the commonly used interfacial modification including physical and chemical treatment on NF, this study provides a design strategy for hierarchical interfaces of polymer composite to prepare high-performance composites with strong interfacial adhesion, featuring simple manufacturing and no toxic reagents.

adhesion and thus is favorable for load transfer for reinforcing the fiber along the interface.59,60 Accordingly, it can be sufficiently concluded that the interfacial properties of the NF/PP composite with interlocking interface were perfected significantly. Further, SEM examinations were conducted to reveal the underlying mechanism of the interfacial enhancement promoted by the interfacially grown NA fibers. The morphologies of the fiber and polymer matrix after the pull-out test are shown in Figure 6. For conventional NF/PP composite, the NF and polymer matrix were characterized by smooth ravine surfaces, without any traces of good interfacial adhesion (Figure 6b,c). The wrinkles distributed along the longitudinal direction stem from the grooves grown naturally at the surface of the pristine NF.21 This provides evidence for weak interfacial adhesion between NF and the polymer matrix. On the contrary, for the NF/PP composite with the interfacially grown NA fibers, some polymers were coated at the surface of NF, an indicator of favorable interfacial adhesion. Moreover, two interfacial failure features were observed in Figure 6d−f: (1) Some NAs granules, which were generated via the fragmentation of NA fibers during the single fiber pull-out test, were adhered onto the NF surface, indicating there is a strong interfacial adhesion between the NF and the grown NA fiber; otherwise, the NA fibers may be directly peeled off the NF and exist as intact fibers, totally different from what were observed in our experiment. This adhesion can be attributed to the strong hydrogen bonding between those two components, as inferred from the shift of the characteristic peak at 3415 cm−1 in the FTIR spectra of Figure 2. Obviously, the interfacial debonding happens until the interfacially grown fibers break. (2) By investigating the polymer matrix as shown in Figure 6e, the microsized slits with some fibrillary ligaments were generated in the matrix, implying that the PP matrix was destroyed with enormous energy dissipation during the debonding processing. Combining the mechanical results and fracture morphologies, a potential interfacial strengthening mechanism is proposed in Figure 7. Ascribing to the NA−NF hydrogen bonding and the transcrystalline layers grown at the NA surface, the interfacially grown NA fibers can bridge NF and PP together to provide remarkable resistance to the interfacial debonding. During the NF pull-out process, the highest stress concentration at the

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CONCLUSION To achieve the mechanical interlocking between PP and NF, the well-defined interfacial NA fibers were constructed on the NF surface by a simple blending approach, followed by the interfacial diffusion and aggregation driven by hydrogen bonding. Ascribing to the NA−NF hydrogen bonding and the transcrystalline layers grown at the NA surface, the laterally grown NA fibers, which rooted firmly on the NF surfaces and penetrated into the PP matrix, integrated PP with NF together via mechanical interlocking, endowing the composite with a favorable interface, as evidenced by the enhanced interfacial shear strength, interfacial friction, and debonding energy. This result sufficiently proved that modulating the interfacial topology is an effective strategy for improving interfacial property of the composites. Based on the simple preparation procedure and high interfacial enhancing efficiency, this interfacial modification is easily put into practice and thus is a significant step toward expanding the potential applications of NF in the advanced polymer/NF composites.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acssuschemeng.7b02448. Additional details about PLM observation and corresponding figure (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(M.N.) E-mail: [email protected]. Fax: +86-28-85402465. Tel.: +86-28-85405133. ORCID

Min Nie: 0000-0001-8386-7547 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is financed by the National Natural Science Foundation of China (51303114 and 51421061), State Key Laboratory of Polymer Materials Engineering (Grant No. sklpme2016-3-05), the Engineering Research Center of Marine Bioresources Comprehensive Utilization, SOA (MBRCU201601), the Program of Innovative Research Team for Young Scientists of Sichuan Province (2016TD0010), and

Figure 7. Interfacial strengthening mechanism of the fiber reinforced composites with the interfacially grown NA fibers during single fiber pull-out test. 10418

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(20) Panapitiya, N. P.; Wijenayake, S. N.; Nguyen, D. D.; Huang, Y.; Musselman, I. H.; Balkus, K. J.; Ferraris, J. P. Gas Separation membranes derived from high-performance immiscible polymer blends compatibilized with small molecules. ACS Appl. Mater. Interfaces 2015, 7, 18618−18627. (21) Xie, L.; Shan, B.; Sun, X.; Tian, Y.; Xie, H.; He, M.; Xiong, Y.; Zheng, Q. Natural fiber-anchored few-layer graphene oxide nanosheets for ultrastrong interfaces in poly(lactic acid). ACS Sustainable Chem. Eng. 2017, 5, 3279−3289. (22) Cao, Y.; Zhang, J.; Feng, J.; Wu, P. Compatibilization of immiscible polymer blends using graphene oxide sheets. ACS Nano 2011, 5, 5920−5927. (23) Nie, M.; Kalyon, D. M.; Fisher, F. T. Interfacial load transfer in polymer/carbon nanotube nanocomposites with a nanohybrid shish kebab modification. ACS Appl. Mater. Interfaces 2014, 6, 14886− 14893. (24) Jin, X.; Heepe, L.; Strueben, J.; Adelung, R.; Gorb, S. N.; Staubitz, A. Challenges and solutions for joining polymer materials. Macromol. Rapid Commun. 2014, 35, 1551−1570. (25) Demir, B.; Henderson, L. C.; Walsh, T. R. Design rules for enhanced interfacial shear response in functionalized carbon fiber epoxy composites. ACS Appl. Mater. Interfaces 2017, 9, 11846−11857. (26) Qian, H.; Kalinka, G.; Chan, K. L. A.; Kazarian, S. G.; Greenhalgh, E. S.; Bismarck, A.; Shaffer, M. S. P. Mapping local microstructure and mechanical performance around carbon nanotube grafted silica fibres: methodologies for hierarchical composites. Nanoscale 2011, 3, 4759−4767. (27) Jin, X.; Strueben, J.; Heepe, L.; Kovalev, A.; Mishra, Y. K.; Adelung, R.; Gorb, S. N.; Staubitz, A. Joining the un-joinable: adhesion between low surface energy polymers using tetrapodal zno linkers. Adv. Mater. 2012, 24, 5676−5680. (28) Cui, C.; Qian, W.; Zhao, M.; Ding, F.; Jia, X.; Wei, F. High strength composites using interlocking carbon nanotubes in a polyimide matrix. Carbon 2013, 60, 102−108. (29) Kim, W. S.; Yun, I. H.; Lee, J. J.; Jung, H.-T. Evaluation of mechanical interlock effect on adhesion strength of polymer−metal interfaces using micro-patterned surface topography. Int. J. Adhes. Adhes. 2010, 30, 408−417. (30) Liu, W.; Nie, M.; Wang, Q. In-situ microfibrillation of polystyrene(PS)/polybutene-1 (PB-1) composites prepared via melt drawing:morphological evolution and properties. J. Polym. Res. 2014, 21, 1−8. (31) Meng, F.; Zhao, R.; Zhan, Y.; Liu, X. Design of Thorn-like micro/nanofibers: fabrication and controlled morphology for engineered composite materials applications. J. Mater. Chem. 2011, 21, 16385. (32) Zheng, N.; Huang, Y.; Sun, W.; Du, X.; Liu, H.-Y.; Moody, S.; Gao, J.; Mai, Y. W. In-Situ Pull-Off of Zno Nanowire from carbon fiber and improvement of interlaminar toughness of hierarchical ZnO nanowire/carbon fiber hydrid composite laminates. Carbon 2016, 110, 69−78. (33) Tamrakar, S.; An, Q.; Thostenson, E. T.; Rider, A. N.; Haque, B. Z.; Gillespie, J. W. Tailoring interfacial properties by controlling carbon nanotube coating thickness on glass fibers using electrophoretic deposition. ACS Appl. Mater. Interfaces 2016, 8, 1501−1510. (34) Patterson, B. A.; Sodano, H. A. Enhanced interfacial strength and uv shielding of aramid fiber composites through zno nanoparticle sizing. ACS Appl. Mater. Interfaces 2016, 8, 33963−33971. (35) An, Q.; Rider, A. N.; Thostenson, E. T. Hierarchical composite structures prepared by electrophoretic deposition of carbon nanotubes onto glass fibers. ACS Appl. Mater. Interfaces 2013, 5, 2022−2032. (36) Bekyarova, E.; Thostenson, E. T.; Yu, A.; Kim, H.; Gao, J.; Tang, J.; Hahn, H. T.; Chou, T. W.; Itkis, M. E.; Haddon, R. C. Multiscale carbon nanotube−carbon fiber reinforcement for advanced epoxy composites. Langmuir 2007, 23, 3970−3974. (37) Qian, H.; Greenhalgh, E. S.; Shaffer, M. S. P.; Bismarck, A. carbon nanotube-based hierarchical composites: a review. J. Mater. Chem. 2010, 20, 4751−4762.

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REFERENCES

(1) Kumar, S. K.; Benicewicz, B. C.; Vaia, R. A.; Winey, K. I. 50th anniversary perspective: Are polymer nanocomposites practical for applications. Macromolecules 2017, 50, 714−731. (2) Taguet, A.; Cassagnau, P.; Lopez-Cuesta, J. M. Structuration, selective dispersion and compatibilizing effect of (nano) fillers in polymer blends. Prog. Polym. Sci. 2014, 39, 1526−1563. (3) Stone, D. A.; Korley, L. T. J. Bioinspired polymeric nanocomposites. Macromolecules 2010, 43, 9217−9226. (4) Liu, Y.; Kumar, S. Polymer/carbon nanotube nano composite fibers–a review. ACS Appl. Mater. Interfaces 2014, 6, 6069−6087. (5) Ramamoorthy, S. K.; Skrifvars, M.; Persson, A. A review of natural fibers used in biocomposites: plant, animal and regenerated cellulose fibers. Polym. Rev. 2015, 55, 107−162. (6) Faruk, O.; Bledzki, A. K.; Fink, H.-P.; Sain, M. Biocomposites reinforced with natural fibers: 2000−2010. Prog. Polym. Sci. 2012, 37, 1552−1596. (7) Dayo, A. Q.; Gao, B. C.; Wang, J.; Liu, W. B.; Derradji, M.; Shah, A. H.; Babar, A. A. Natural hemp fiber reinforced polybenzoxazine composites: curing behavior, mechanical and thermal properties. Compos. Sci. Technol. 2017, 144, 114−124. (8) Guo, F.; Wang, N.; Cheng, Q.; Hou, L.; Liu, J.; Yu, Y.; Zhao, Y. Low-cost coir fiber composite with integrated strength and toughness. ACS Sustainable Chem. Eng. 2016, 4, 5450−5455. (9) Gardea, F.; Glaz, B.; Riddick, J.; Lagoudas, D. C.; Naraghi, M. Energy dissipation due to interfacial slip in nanocomposites reinforced with aligned carbon nanotubes. ACS Appl. Mater. Interfaces 2015, 7, 9725−9735. (10) Ehlert, G. J.; Sodano, H. A. Zinc Oxide Nanowire Interphase for enhanced interfacial strength in lightweight polymer fiber composites. ACS Appl. Mater. Interfaces 2009, 1, 1827−1833. (11) Tzounis, L.; Kirsten, M.; Simon, F.; Mäder, E.; Stamm, M. The interphase microstructure and electrical properties of glass fibers covalently and non-covalently bonded with multiwall carbon nanotubes. Carbon 2014, 73, 310−324. (12) Muthuraj, R.; Misra, M.; Mohanty, A. K. Injection molded sustainable biocomposites from poly(butylene succinate) Bioplastic and perennial grass. ACS Sustainable Chem. Eng. 2015, 3, 2767−2776. (13) Ning, N. Y.; Fu, S. R.; Zhang, W.; Chen, F.; Wang, K.; Deng, H.; Zhang, Q.; Fu, Q. Realizing the enhancement of interfacial interaction in semicrystalline polymer/filler composites via interfacial crystallization. Prog. Polym. Sci. 2012, 37, 1425−1455. (14) Ma, S.; Ye, Q.; Pei, X.; Wang, D.; Zhou, F. Antifouling on gecko’s feet inspired fibrillar surfaces: evolving from land to marine and from liquid repellency to algae resistance. Adv. Mater. Interfaces 2015, 2, 1500257−1500257. (15) Chen, F.; Liu, W.; Seyed Shahabadi, S. I.; Xu, J.; Lu, X. Sheetlike lignin particles as multifunctional fillers in polypropylene. ACS Sustainable Chem. Eng. 2016, 4, 4997−5004. (16) Wang, Y.; Raman Pillai, S. K.; Che, J.; Chan-Park, M. B. High interlaminar shear strength enhancement of carbon fiber/epoxy composite through fiber- and matrix-anchored carbon nanotube networks. ACS Appl. Mater. Interfaces 2017, 9, 8960−8966. (17) Karmarkar, A.; Chauhan, S. S.; Modak, J. M.; Chanda, M. Mechanical properties of wood−fiber reinforced polypropylene composites: effect of a novel compatibilizer with isocyanate functional group. Composites, Part A 2007, 38, 227−233. (18) Zhang, K.; Misra, M.; Mohanty, A. K. Toughened sustainable green composites from poly(3-hydroxybutyrate-co-3-hydroxyvalerate) based ternary blends and miscanthus biofiber. ACS Sustainable Chem. Eng. 2014, 2, 2345−2354. (19) Johari, A. P.; Mohanty, S.; Kurmvanshi, S. K.; Nayak, S. K. Influence of different treated cellulose fibers on the mechanical and thermal properties of poly(lactic acid). ACS Sustainable Chem. Eng. 2016, 4 (3), 1619−1629. 10419

DOI: 10.1021/acssuschemeng.7b02448 ACS Sustainable Chem. Eng. 2017, 5, 10413−10420

Research Article

ACS Sustainable Chemistry & Engineering (38) Karger-Kocsis, J.; Mahmood, H.; Pegoretti, A. Recent advances in fiber/matrix interphase engineering for polymer composites. Prog. Mater. Sci. 2015, 73, 1−43. (39) He, X.; Li, Y.; Nie, M.; Wang, Q. Root-like glass fiber with branched fiber prepared via molecular self-assembly. RSC Adv. 2016, 6, 45492−45494. (40) He, X.; Xu, W.; Liu, Y.; Nie, M.; Wang, Q. In situ construction of pompon-like hydroxyapatite hybrid via interfacial self-assembly in polypropylene matrix. Compos. Sci. Technol. 2017, 142, 246−252. (41) Wang, G. W.; Chen, B.; Zhuang, L. H.; Yun, K.; Guo, J. R.; Wang, Y.; Xu, B. Dyeing performances of ramie fabrics modified with an amino-terminated aliphatic hyperbranched polymer. Cellulose 2015, 22, 1401−1414. (42) Luo, F.; Geng, C.; Wang, K.; Deng, H.; Chen, F.; Fu, Q.; Na, B. New understanding in tuning toughness of β-polypropylene: the role of β-nucleated crystalline morphology. Macromolecules 2009, 42, 9325−9331. (43) Zhou, T.; Zhang, A.; Zhao, C. S.; Liang, H. W.; Wu, Z. Y.; Xia, J. K. Molecular chain movements and transitions of SEBS above room temperature studied by moving-window two-dimensional correlation infrared spectroscopy. Macromolecules 2007, 40, 9009−9017. (44) Hu, D.; Wang, G.; Feng, J.; Lu, X. Exploring supramolecular self-assembly of a bisamide nucleating agent in polypropylene melt: the roles of hydrogen bond and molecular conformation. Polymer 2016, 93, 123−131. (45) Shih, R. S.; Lu, C. H.; Kuo, S. W.; Chang, F. C. Hydrogen bondmediated self-assembly of polyhedral oligomeric silsesquioxane-based supramolecules. J. Phys. Chem. C 2010, 114, 12855−12862. (46) Noda, I. Generalized Two-dimensional correlation method applicable to infrared, raman, and other types of spectroscopy. Appl. Spectrosc. 1993, 47, 1329−1336. (47) Noda, I. Determination of two-dimensional correlation spectra using the hilbert transform. Appl. Spectrosc. 2000, 54, 994−999. (48) Watanabe, A.; Morita, S.; Ozaki, Y. Study on temperaturedependent changes in hydrogen bonds in cellulose Iβ by infrared spectroscopy with perturbation-correlation moving-window twodimensional correlation spectroscopy. Biomacromolecules 2006, 7, 3164−3170. (49) Xue, B.; Zhang, J.; Zhou, T. Moving-window two-dimensional correlation infrared spectroscopic study on the dissolution process of poly(vinyl alcohol). Anal. Bioanal. Chem. 2015, 407, 8765−8771. (50) Guo, Y.; Wu, P. Investigation of the hydrogen-bond structure of cellulose diacetate by two-dimensional infrared correlation spectroscopy. Carbohydr. Polym. 2008, 74, 509−513. (51) Peng, L.; Zhou, T.; Huang, Y.; Jiang, L.; Dan, Y. Microdynamics mechanism of thermal-induced hydrogel network destruction of poly(vinyl alcohol) in D2O studied by two-dimensional infrared correlation spectroscopy. J. Phys. Chem. B 2014, 118, 9496−9506. (52) Deshmukh, Y. S.; Wilsens, C. H. R. M.; Leoné, N.; Portale, G.; Harings, J. A. W.; Rastogi, S. Melt-miscible oxalamide based nucleating agents and their nucleation efficiency in isotactic polypropylene. Ind. Eng. Chem. Res. 2016, 55, 11756−11766. (53) Han, R.; Li, Y.; Wang, Q.; Nie, M. Critical formation conditions for β-form hybrid shish-kebab and its structural analysis. RSC Adv. 2014, 4, 65035−65043. (54) Menczel, J.; Varga, J. Influence of nucleating agents on crystallization of polypropylene: i. talc as a nucleating agent. J. Therm. Anal. 1983, 28, 161−174. (55) Abdou, J. P.; Braggin, G. A.; Luo, Y.; Stevenson, A. R.; Chun, D.; Zhang, S. Graphene-induced oriented interfacial microstructures in single fiber polymer composites. ACS Appl. Mater. Interfaces 2015, 7, 13620−13626. (56) Sharma, M.; Gao, S.; Mäder, E.; Sharma, H.; Wei, L. Y.; Bijwe, J. Carbon fiber surfaces and composite interphases. Compos. Sci. Technol. 2014, 102, 35−50. (57) Hampe, A.; Kalinka, G.; Meretz, S.; Schulz, E. An advanced equipment for single-fibre pull-out test designed to monitor the fracture process. Composites 1995, 26, 40−46.

(58) Zhandarov, S.; Mäder, E. Characterization of fiber/matrix interface strength: applicability of different tests, approaches and parameters. Compos. Sci. Technol. 2005, 65, 149−160. (59) Graupner, N.; Rößler, J.; Ziegmann, G.; Müssig, J. Fibre/matrix adhesion of cellulose fibres in pla, pp and mapp: a critical review of pull-out test, microbond test and single fibre fragmentation test results. Composites, Part A 2014, 63, 133−148. (60) Thomason, J. L.; Vlug, M. A. Influence of fibre length and concentration on the properties of glass fibre-reinforced polypropylene: 4. impact properties. Composites, Part A 1997, 28, 277−288.

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DOI: 10.1021/acssuschemeng.7b02448 ACS Sustainable Chem. Eng. 2017, 5, 10413−10420