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Nanofibril Organization in Silk Fiber as Inspiration for Ductile and Damage-Tolerant Fiber Design

  • Shihui Lin
  • Chao Ye
  • Wenwen Zhang
  • Anchang XuEmail author
  • Shixian Chen
  • Jing RenEmail author
  • Shengjie LingEmail author
Letter
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Abstract

Ductile and damage-tolerant fibers (DDTFs) are desired in aerospace engineering, mechanical engineering, and biomedical engineering because of their ability to prevent the catastrophic sudden structural/mechanical failure. However, in practice, design and fabrication of DDTFs remain a major challenge due to finite fiber size and limited processing techniques. In this regard, animal silks can provide inspirations. They are hierarchically structured protein fibers with an elegant trade-off of mechanical strength, extensibility and damage tolerance, making them one of the toughest materials known. In this article, we confirmed that nanofibril organization could improve the ductility and damage-tolerance of silk fibers through restricted fibril shearing, controlled slippage and cleavage. Inspired by these strategies, we further established a rational strategy to produce polyamide DDTFs by combining electrospinning and yarn-spinning techniques. The resultant polymeric DDTFs show a silk-like fracture resistance behavior, indicating potential applications in smart textile, biomedicine, and mechanical engineering.

Graphic Abstract

Keywords

Fiber Silk Mechanical property Damage tolerance Bioinspired materials 

Introduction

The sudden accidental rupture happens in brittle fibers is of great concern for practical applications [1]. Ductile and damage-tolerant fibers (DDTFs), hence, are desired in many fields, such as aerospace engineering, civil engineering, and biomedical engineering [2, 3, 4]. DDTFs have a strong ability to sustain the main structure in the defect area safely before catastrophic damage happens [2]. According to Ludwik-Davidenkov-Orowan theory [5], whether a material is brittle or ductile depends on the relative value of its yield stress and fracture stress. If the fracture stress is much lower than the yield stress, brittle fracture occurs. By contrast, if the fracture stress is much higher than the yield stress, a yield point can be observed, and thereby the ductile fracture appears. As a result, the ductile fibers feature a typical J-shaped stress–strain curve with an elastic phase (from 0 to yield point) and a plastic deformation region (after yield point) [3, 6]. Damage tolerance, on the other hand, is a property of a structure relating to its ability to absorb energy when a crack grows, which plays a critical role in preventing the case of sudden catastrophic failure [1].

In engineering material designs, the ductile and damage-tolerant materials can be achieved by creating metastructures [7, 8], designing nanocomposites [9, 10, 11, 12, 13] or directly printing layered/patterned 3D architectures [14, 15, 16, 17]. However, these employed techniques, such as laser cutting, vacuum filtration, self-assembly, or 3D-printing, require the materials to have large enough sizes or to withstand violent processing. These extreme conditions hindered the applications of strategies in which hinders the development in DDTFs fabrication. By far, it remains a significant challenge to produce DDTFs by fiber engineering. However, efficient strategies can be found in nature [18, 19]. Wild silkworm cocoon silk as an example, despite the presence of defects, such as cavities and tears, it still displays remarkable mechanical properties with a unique combination of high fracture strength (300–700 MPa) and extensibility (30–40%) [20, 21]. Such mechanical feature enables silk to absorb a large amount of energy before breaking, results in toughness values that are several times higher than high-performance steel and Kevlar fibers [19].

Computational modeling has suggested the nanofibrils play an essential role in governing the fracture resistance of silk fibers, but sufficient experimental proof is still needed [20, 22, 23, 24]. Accordingly, we investigate the contribution of nanofibrils on fracture mechanics of the natural silk fibers through the microscopic camera system cooperated with the in situ tensile test. Our findings confirm that nanofibrils can improve the ductility and damage tolerance of the silk fibers through non-slip kinematics, restricted fibril shearing, and controlled slippage. Learning from these natural designs, we further develop silk fiber inspired DDTFs by using polyamide as starting materials. Electrospinning and wet-spinning are combined to spin continuous long fibers with nanofibrous structures. The resultant polymeric DDTFs feature silk fiber-like ductile and damage-tolerant behaviors, showing potential applications in functional textile, biomedicine, and mechanical engineering.

Experimental Section

Preparation of Silk Fibers

All silk fibers used in this work are force-reeled Antheraea pernyi (A. pernyi) silk fibers [25]. A. pernyi silkworms were obtained from the wild tussah forest in Dandong, Liaoning province, China. When silkworms start spinning, the head of the silkworm was fixed to prevent it from swinging from side to side. The silk fibers were directly pulled out from the mouthpart of the silkworm with controlled speed (20 mm/s) and collected by a rotating cylinder.

Preparation of Polyamide DDTFs

Polyamide 6 in formic acid with a concentration of 25 wt% was firstly prepared at room temperature. The coagulation bath consists of 0.05 wt% Peregal O aqueous solution with a copper sheet at the bottom of the container. A 15 mL syringe filled with the resultant polymer solution was mounted on a syringe pump, and a needle with an internal diameter of 0.8 mm was used. The syringe pump can provide constant solution droplets at the tip of the needle at the desired speed of 0.2 mL/h. An electric potential of 20 kV was applied between the needle and the coagulation bath from a high-voltage–power supply. During the electrospinning, the as-spun nanofibers were firstly immersed into the coagulation bath. A yarn spinning processing was then carried out to obtain nanofibrous DDTFs. At last, the fabricated DDTFs were passed through a water bath and collected by an automatic roller.

Notching on Silk Fibers

We introduced artificial notches in two ways. For the samples used to observe the fracture morphology, we used commercial laser microdissection (Leica LMD 7, Wetzlar, Germany) to make a smooth notch on silk fibers. Specifically, a single fiber was first mounted on the self-cutting cardboard frame with the base length of about 10 mm and fixed with cyanoacrylate. Then, the resultant sample was mounted on the objective table of the laser microdissection. At last, a laser with a power of 55 mW and wavelength of 349 nm was used to cut the selected area of the samples, leaving artificial notches on the silk fibers. In these experiments, artificial notches with widths of ~5 to 20 μm and depth of ~ 2 to 10 μm was introduced in the surface of the single silk fibers. For the samples used for testing mechanical behaviors, we used blades to notch the surface of samples directly under a stereoscope, which can prevent the effects caused by laser damage. The artificial notches were in varying depth, but all of the depths were less than half of the sample diameter.

Mechanical Tests

Sample preparation: the sample preparations of both silk fibers and polyamide DDTFs were the same. A single fiber was mounted on the self-cutting cardboard frame with the base length of about 20 mm and fixed with cyanoacrylate. About 15 mm of adjacent fibers were reserved for calculating the cross-sectional area.

Tensile measurements: After overnight drying of cyanoacrylate, the samples were installed on the test machine (Instron 5966 machine, Instron, Norwood, USA). Both sides of the board frames were cut before the test, and the height of the frame was adjusted to maintain the fibers at zero load point, and then measured the initial length of the fiber with calipers. The tensile tests of silk fibers were conducted at room temperature, with a relative humidity of 75% and a tensile speed of 2 mm/min. The tensile tests of polyamide DDTFs and commercial polyamide fibers were both conducted at room temperature with a relative humidity of 70% and a tensile speed of 2 mm/min.

X-ray Diffraction Experiments

Small-angle X-ray scattering (SAXS) was performed to investigate the orientation of silk fibers and DDTFs. SAXS experiments were carried out at Characterization and Analysis Center of ShanghaiTech University by using Xenocs equipment, Xeuss 2.0. The diffraction patterns were collected by the detector with 172 pixels × 172 pixels of 172 μm × 172 μm area each. The wavelength and the photon flux of the X-ray source were 1.54189 Å, 4.0 × 107 photons s−1, respectively. The beam size at the detector was 1.2 mm × 1.2 mm. X-ray diffraction (XRD) curve of polyamide DDTFs was recorded on Bruker D8 Advance diffractometer with Cu Kα radiation with the voltage of 40 kV. Data were collected in the 2θ range of 10°–35° with the steps of 0.02°.

Density Measurement of Polyamide DDTFs

We took a certain length of DDTF (more than 1 m) and weighed its mass. The density of DDTFs was calculated by dividing mass by volume. The volume was calculated by multiplying the average cross-sectional area by the length.

Characterization

The cross-sectional areas of the silk fibers and polyamide DDTFs were observed by scanning electron microscope (SEM, Phenom Pro). SEM images were further analyzed by the ImageJ software to estimate the cross-sectional area of each fiber. The cross-sectional areas of both notched fibers and unnotched fibers were calculated according to the intact part of the fiber. The average value was used for the stress calculation.

The surface and cross-sectional morphologies of silk fibers and polyamide DDTFs were observed by high-resolution SEM (JEOL JSM-7800F, Tokyo, Japan) at an acceleration voltage of 5 kV. All the samples were coated with a 5 nm thick gold layer to provide conductivity before observation.

Results

Nanofibril Organization in Silk Fibers

Figure 1a depicts the typical structural of the silk fiber, which can be considered as a coaxial fiber composed of a core and a shell layer [26, 27, 28]. The shell layer, also called sericin, constitutes about 10–30% weight of the whole fiber but does not essentially contribute to the tensile properties [25]. The highly organized core silk fibroin filaments are the primary source to provide the fiber’s mechanical properties. These silk filaments are composed of microfibrils with several hundred nanometers in diameter (Fig. 1b). At the smaller scale, these microfibrils are constituted of highly oriented nanofibrils (about 5–200 nm in diameter) that are composed of such as β-sheet, random coil, and helical structure. SAXS pattern of the force-reeled A. pernyi silks is shown in Figure S1. According to a reciprocal relation between the orientation of scatterer and the scattering patterns, the vertical scattering patterns represent the minor axis direction of micro slits of the fibers, and the horizontal scattering patterns reflect the long axis direction of micro slits of the fibers. The SAXS pattern indicates that the microfibrils in force-reeled A. pernyi silks are highly oriented. This sophisticated hierarchical structure makes the silk fibers much stronger and tougher than most of the natural and synthetic polymer fibers, such as camel hair, polyamide and polyurethane fibers (Fig. 1c). Ashby plot in Fig. 1d further reveals that silk fibers are better at balancing strength and stiffness than other natural and synthetic materials [29].
Fig. 1

Hierarchical structure and mechanical performance of force-reeled A. pernyi silk fibers. a Schematic of the hierarchical structure of force-reeled A. pernyi silk fibers. b Typical SEM image of the cross-section of force-reeled A. pernyi silk fiber. c Typical stress–strain curves of A. pernyi silk fiber, camel hair, polyamide (PA) and polyurethane (PU) fiber. d Ashby plot compares the specific strength and specific stiffness of force-reeled A. pernyi silk fiber with other natural and synthetic materials. Ashby plot of natural and synthetic materials are adapted from Ref. [29]

The Contribution of Nanofibril on Fracture Mechanics of Single Silk Fiber

In order to examine the contribution of nanofibrils on crack propagation of the silk fibers, we integrated a tensile test with a microscopic camera system to monitor the tensile failure process of the single notched silk fiber. In these experiments, artificial notch (with widths of ~ 5 to 20 μm and depth of ~ 2 to 5 μm) was introduced in the surface of the single silk fiber, and the mechanical properties and fracture process were tested to compare with that of the adjacent intact (unnotched) fiber (Fig. 2a, b). Compared with silk fiber extracted from silkworm cocoons, the force-reeling silks used in this work, which are directly reeled from silkworm spinneret, have better mechanical properties and structural uniformity. Because the natural spinning process of silkworm usually introduces defects when the silkworm swings the head. Therefore, the force-reeling silk without these natural defects has proven to be a better candidate to study the structure–property relationships. Figure 2c shows the typical stress–strain curves of notched and adjacent unnotched single silk fibers. When the notch width is smaller than the half-width of the fiber, the notched fibers exhibited the similar stress–strain curves as the unnotched fibers; only the failure to strain was reduced, exhibiting a typical ductile fracture behavior.
Fig. 2

Fracture mechanics of single force-reeled A. pernyi silk fiber. a Illustration of the sample preparation method. b Optical microscope image of force-reeled A. pernyi silk fiber with an artificial incision. c Typical stress–strain curves of notched and unnotched force-reeled A. pernyi fiber. The cross-sectional area of the notched silk fiber was calculated according to the intact part of the fiber. d Fracture process of a notched force-reeled A. pernyi silk fiber during the tensile test. The images presented are extracted from the video frames of Movie S1. The Movie was recorded by optical microscopy. The test was conducted at room temperature with a tensile speed of 2 mm/min and relative humidity of 75%. The tensile directions were indicated by the white hollow arrows. e, f SEM image of notched force-reeled A. pernyi silk fiber after tensile failure. f The high-resolution SEM image of the cleaved area of the silk fiber in e. g Schematic diagram of the theoretical model of Griffith’s size scaling. The schematic is adapted from ref [4], copyright 2013 John Wiley and Sons

As shown in Fig. 2d–f, the ductile behavior of silk fibers is achieved by the restricted fibril shearing and controlled slippage (Fig. 2d) and cleavage (Fig. 2e, f), which is also supported by the previous simulation findings [20, 24]. The microscopic camera snapshots and the SEM observation further confirm the damage tolerant behavior of single silk fibers. During the tensile processing, the initial crack was deflected around 90 degrees (the direction indicated by the white arrows), and the apparent nanofibril shearing, splitting and pulling were observed.

As demonstrated in Fig. 2g, nanofibril can further improve the failure strength and toughness through the increase of the size of the process zone [4]. According to Griffith’s size scaling, the failure strength σf of a polymer fiber with diameter of D and free of intrinsic flaws can reach the theoretical limit of the interatomic bonds, denoted by σth, whereas a fiber containing a large flaw significantly decreases its strength, which can be quantitatively expressed as follows:
$${\upsigma }_{f} \sim{\raise0.7ex\hbox{$1$} \!\mathord{\left/ {\vphantom {1 {\sqrt D }}}\right.\kern-0pt} \!\lower0.7ex\hbox{${\sqrt D }$}}$$
(1)
For the notched silk fibers, the failure strength σf (~ 250 MPa, as shown in Fig. 2c) is approximately equal to theoretical strength σth. Herein, σth can be assessed according to the following empirical equation [28]:
$${\upsigma }_{th} \approx \sqrt {{\raise0.7ex\hbox{${12\varGamma {\text{E}}}$} \!\mathord{\left/ {\vphantom {{12\varGamma {\text{E}}} d}}\right.\kern-0pt} \!\lower0.7ex\hbox{$d$}}}$$
(2)
where Γ is surface energy, E is the modulus, d is the size of the characteristic structure in silk fibers. Specifically, a silk fiber with nanofibril size of 120 nm (evaluated from SEM images in Fig. 2f) and a modulus of 6 GPa with a surface energy of 0.15 J m−2 would be expected to have σth of about ~ 300 MPa.

The condition that σf ≈  σth can be explained by the process zone, also referred to as “cavitation box,” which characterizes the number of nanofibrils that contributes to resisting fracture (Fig. 2g) [4]. Compared with the isotropic fibers without nanostructure, nanofibrils in silk fiber significantly increase the size of the process zone, with a value of l0 ≈ D, driving a crack through the fiber leads to widespread damage that is not only limited to the crack surface. Furthermore, in this case, the stress concentrations at the crack tip are diminished, which can further reduce the threat imposed by cracks of leading to catastrophic damage.

Silk Fiber Inspired Polyamide DDTFs

Inspired by such a natural nanofibril organization, we further used electrospinning to produce a silk fiber inspired DDTFs based on polyamide starting materials. Electrospinning has been registerd as a facile fabrication technique, however, it's typical model is unable to spin nanofibers in a continuous and oriented form. So, here we integrated the electrospinning and wet-spinning techniques to achieve this goal. As shown in Fig. 3a, this combined spinning system comprises three primary systems: an electrospinning device, a water bath, and a rotating drum. During the process, the as-spun polymer nanofibers, produced from electrospinning device, are first coagulated into nanofiber mats on the surface of the water bath. Continuous post-rolling processing is then conducted on nanofiber mats. This process helps the nanofiber mats gather and form a single fiber, in which the nanofibers align along with the drawing direction (Fig. 3a).
Fig. 3

Preparation and structural characterization of polyamide DDTFs. a Schematic of the strategy for producing polyamide DDTFs. In this strategy, electrospinning and yarn spinning method are combined to prepare continuously spinnable polyamide DDTFs. b Photograph of polyamide DDTFs spun from 25 wt% polyamide 6-formic acid solution. c SEM image of the surface of polyamide DDTFs. d SEM image of the notched polyamide DDTFs. The notch was created by a laser microdissection system

Using this strategy, we can continuously collect the highly oriented nanofibrous fibers with a length of more than 20 km at the rolling speed of 4 m/min. The resultant polyamide DDTFs has a cross sectional area of 2446 ± 627 μm2 with 0.4 g cm–3 in density (Fig. 3b). The uniformity of the produced fibers is influenced by the collection process. Fibers are stretched when under the high reeling speed, resulting in fiber width varies from 62 to 276 μm. The uniformity can be improved by post-stretching or matching each components in the collection part to reduce uneven stress.

In this study, single needle based electrospinning device was used to produce polyamide DDTFs. In this way, we usually collect 90 m of polyamide DDTFs in 1 h of spinning (with a reeling speed of 2.5 cm s–1). Such a production efficiency is inferior to that of commercialized wet-spinning techniques, which often can spin more than 3000 m of polyamide fibers in an hour. The production efficiency of electrospinning can be improved by using multi-needle based strategy. However, the major challenge of multi-needle based technique is electric field interferences between the needles, which has a significant effect on the structural uniformity and mechanical performance of the resultant fibers.

Figure 3c shows the typical surface morphologies of the polyamide DDTFs. In contrast to the smooth surface of commercialized polyamide fibers, the polyamide DDTFs show highly oriented nanofibrous structures that formed by the as-spun nanofibers. The high-resolution SEM image reveals that these nanofibers fuse together and align along the fiber axis with significant weak interfaces between nanofibers, sharing silk fiber-like nanostructures. In addition, these nanofibers have high uniformity with an average width of 1.9 ± 0.3 μm. The SAXS pattern of polyamide DDTFs (Figure S2) also agrees well with orientation structure of the fibers. XRD experiment was carried out to determine the crystallinity of polyamide DDTFs (Figure S3). The crystallinity xc is given by
$$x_{c} \left( \% \right) = \frac{{A_{\text{c}} }}{{A_{\text{a}} + A_{\text{c}} }} \times 100$$
(3)
where Aa is the area due to amorphous diffusion (shaded area in Figure S3); Ac is the area of crystalline peaks. The calculating crystallinity of polyamide DDTFs is about 25%.
Next, we focus on testing the fracture mechanics of these DDTFs. As shown in Fig. 3d, before the tensile measurement, a notch with a width of one-third of the original fiber width has been created on the polyamide DDTFs. As shown in Fig. 4a-c, similar with natural silk fiber, the notched polyamide DDTFs exhibit excellent damage-tolerant behavior. Characteristic ductile fractures are detected from stress–strain plots (Fig. 4a). Typical crack shifting and nanofiber pulling are also found during the tensile processing. The crack propagation is prevented by these highly oriented nanofibers (Fig. 4b, c, Movie S2). By contrast, the stress–strain curve of notched commercial polyamide (nylon 66) fiber, a fiber without the nanofibril structure (Fig. 4c, d), shows an entirely different fracture behavior with the unnotched one (Fig. 4d, e). Linear crack propagation (Fig. 4f) is detected in the notched fiber, which fractures before reaching the yield point (a strain of ~ 10%). A similar phenomenon can also be seen in notched commercial polypropylene (PP) fiber (Figure S3). The mechanical properties of PP fiber, polyamide fiber, and the polyamide DDTF before and after introducing notches are further compared (Table S1). The polyamide DDTF shows significantly better strain retention of 80.81% (17.08% and 20.85% for PP fiber and polyamide fiber) and improved stress retention of 39.56% compared to 28.71% of polyamide fiber.
Fig. 4

Mechanical properties of polyamide DDTFs and commercial polyamide fibers. a Stress–strain profile of notched and unnotched polyamide DDTFs (the sectional areas of the notched polyamide DDTFs were calculated according to the intact part of each fiber and the test was conducted at room temperature with a tensile speed of 2 mm/min and relative humidity of 70%). b Illustration of the fracture process of polyamide DDTFs that observed from high-speed camera system (Movie S2). During tensile failure, a notched polyamide DDTFs undergo crack shifting and nanofiber pulling. c SEM image of the fracture surface of the polyamide DDTFs. d SEM image of the cross-section of the commercial polyamide fiber. e Stress–strain curves of notched and unnotched commercial polyamide fibers (the sectional areas of the notched commercial polyamide fibers were calculated according to the intact part of each fiber and the test was conducted at room temperature with a tensile speed of 2 mm/min and relative humidity of 70%). f Fracture process of notched commercial polyamide fiber. Rapid crack propagation is observed

Discussion

Of note, as shown in Fig. 4a, the tensile strength of DDTFs is almost ten times lower than the nylon 66 fiber. This appearance agrees well with the fact that many manmade materials bear a conflict between mechanical strength and ductility [3]. Strong materials, such as high-performance Kevlar and carbon fiber are often brittle, while elastic fibers, such as polyurethane fibers, instead, are often weak [19]. However, this conflict does not exist in silk fibers, especially the spider dragline silk and wild silkworm cocoon silks, which feature a unique combination of strength and mechanical toughness. These animal silk fibers resolve the conflict through hierarchical and heterogeneous nanoconfinement, which is defined as optimizing the strength, stiffness, and toughness of materials through confining their building blocks at critical length scales [30, 31].

In animal silks, nanoconfinement exists in multiple length scales (Fig. 5). At the nanocrystal scale, the β-sheet structures are confined at 3–4 nm in length. These β-sheets serve as cross-linkers to connect the amorphous regions into network structures. The high strength of β-sheets originates from nanoconfinement and results in cooperation of hydrogen bonds against shear. The interaction between β-sheets and amorphous regions leads to a nonlinear stiffening behavior of silk [30]. At the silk nanofibril scale, the nanoconfinement of silk nanofibril size leads to homogeneous deformation of all silk proteins in the silk nanofibrils, achieving optimized strength and toughness through contributions of many protein components (e.g., β-sheets, β-turns, helices, and random coils) acting together. At a higher level of the hierarchy, silk nanofibrils with several geometrical confinements (e.g., confined diameters and heterogeneous globular structures) are bundled into fibers to enhance mechanical properties further. All these ingenious hierarchical organizations contribute to making spider silk extremely strong, resilient, and robust against defects [23, 32].
Fig. 5

Multiscale nanoconfinement in animal silks. The nanoconfinement exists in crystal size, protein composite, and nanofibril dimension, which makes the animal silks ductile and damage tolerance

For bioinspired polyamide DDTFs, currently, we can only mimic the nanofibril organization of silk fibers. Challenge remains on mimicking their more refined natural nanostructures, such as nanocrystal confinement and the nanocrystal-amorphous interfaces, which also play the critical roles in governing the strength and toughness of the silk fiber. This gap between natural silk and bioinspired DDTFs prevents us from generating polymeric DDTFs with a comprehensive mechanical property that comparable with the silk fibers. However, this study reveals that the simplified natural models still can inspire the engineering material design and fabrication, especially, for the researches which aim to obtain a material with customized mechanical requirements. Besides, natural fibrils composed of cellulose, chitin, and silk nanofibrils are excellent engineering materials, which not only endow the exceptional mechanical properties of biological structural materials but also support physiological functions. Smart formation/assembling principles provide outstanding properties, such as nanoconfinement and nanofibril orientation usually contribute a unique combination of high tensile strength and extensibility; 3D helicoidally stacking produces high resist fracture efficiently through energy dissipation and prohibit crack propagation. Thus, further research is desired to take advantage of natural materials to achieve superior artificial fiber materials [33, 34, 35, 36]. Finally, some functional fiber devices, such as fiber sensor and conformal or deformable smart textiles, require the fiber with high flexibility and fracture resistance rather than the high mechanical strength or modulus.

Conclusions

In summary, our experimental results reveal that nanofibril organization in silk fibers plays a vital role in channeling the ductility and damage-tolerance of silk fibers, achieving by the non-slip kinematics, restricted fibril shearing, and controlled slippage. Learning from nanofibril organization in natural silk fibers, we further fabricate synthetic polymer-sourced DDTFs through the integrated electrospinning and yarn-spinning techniques. Compared with commercial polymer fibers with identical chemical constitutions, the silk fiber-inspired polyamide DDTFs feature highly oriented nanofibrous structures, which channel the ductility and damage-tolerance of the fibers in the same manner as observed in silk fiber. These DDTFs show promising applications in smart textile, biomedicine, and mechanical engineering. Furthermore, our strategies in producing nanofibrous polymer can be applied for different polymer spinning fiber systems leading to enhanced  mechanical and functional properties.

Notes

Acknowledgements

We acknowledge National Natural Science Foundation (No. 51973116, U1832109, 21935002), Shanghai Pujiang Program (18PJ1408600), the National Natural Science Foundation of China (21808220), the starting grant of ShanghaiTech University and Shanghai Sailing Program (17YF1411500) for support of this work.

Compliance with Ethical Standards

Conflict of interest

The authors declare no competing interests.

Supplementary material

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Supplementary material 1 (DOCX 1271 kb)

Supplementary material 2 (MP4 11686 kb)

Supplementary material 3 (MP4 3673 kb)

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Copyright information

© Donghua University, Shanghai, China 2019

Authors and Affiliations

  1. 1.School of Physical Science and TechnologyShanghaiTech UniversityShanghaiChina
  2. 2.Jiangsu Co-Innovation Center of Efficient Processing and Utilization of Forest Resources, Jiangsu Key Lab of Biomass-Based Green Fuel and Chemicals, College of Chemical EngineeringNanjing Forestry UniversityNanjingChina
  3. 3.State Key Laboratory of New Textile Materials and Advanced Processing Technology, School of Textile Science and EngineeringWuhan Textile UniversityWuhanChina

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