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Poly(l -lactic acid) twisted nanofiber yarn prepared by carbon dioxide laser supersonic multi-drawing
European Polymer Journal 110 (2019) 145–154
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Poly(L-lactic acid) twisted nanofiber yarn prepared by carbon dioxide laser supersonic multi-drawing
T
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Akihiro Suzuki , Yuta Shimba University of Yamanashi, Interdisciplinary Graduate of School, Takeda-4, Kofu 400-8511, Japan
A B S T R A C T
Poly(L-lactic acid) (PLLA) twisted nanofiber yarn (TNFY) was continuously prepared using a CO2-laser supersonic multi-drawing apparatus equipped with a twisterwinder that offers independent control of winding speed (SW) and twisting speed (ST). TNFYs with various twist numbers (nT), twist angles, and bulk densities were produced by changing SW and ST. TNFYs with a wide nT range of 395–1276 T m−1 were obtained. Furthermore, the mechanical properties of the obtained TNFYs depended upon the degree of twist. The maximum Young’s modulus was 3.06 cN dtex−1 at ST = 300 T min−1 and SW = 0.440 m min−1, and the maximum tensile strength was 0.36 cN dtex−1 at ST = 365 T min−1 and SW = 0.440 m min−1. The Young’s modulus had a maximum value at optimum obliquity because the high twist reduced the contribution of fiber strength to the yarn hardness due to fiber obliquity in the yarn. The tensile strength increased with increasing ST because the frictional forces among the NFs were increased by twisting at higher ST.
1. Introduction Nanofibers (NFs) are used in a variety of applications, such as membranes [1–3], biomedical devices [4], and scaffolds for tissue engineering [5–9]. NFs have been produced by electrospinning [10–21], sea-island-type conjugated melt spinning, and melt blowing [22,23]. We previously proposed a technique for the production of NFs, referred to as a carbon dioxide (CO2) laser supersonic drawing (CLSD). It involves irradiating a fiber with radiation from CO2-laser while drawing it at a supersonic velocity. A supersonic jet was generated by blowing air into a vacuum chamber through the orifice used to inject the fiber into the vacuum chamber. The adiabatic expansion of air across the orifice cools the jet. The fiber is instantly melted by the high-power laser beam that irradiates in the cold supersonic jet. It is then tremendously deformed by the shear force generated by the supersonic flow, and ultradrawn to a draw ratio of the order of 105. CLSD is a novel technique that uses laser irradiation of fibers in a supersonic jet, and can easily prepare NFs of various polymers. CLSD employs only CO2 laser irradiation and does not require additional processes or solvents, unlike electrospinning, which requires a solvent, or sea-island-type conjugated melt spinning that requires the removal of a second component. CLSD has already been applied to prepare NFs of poly(L-lactic acid) (PLLA) [24], poly(ethylene terephthalate) [25], poly (ethylene-2,6-naphthalate) [26], poly(glycolic acid) [27], ethylene tetrafluoroethylene copolymer [28], nylon 66 [29], and poly(phenylene sulfide) [30]. Modification of the CLSD technique has led to the development of
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CO2-laser supersonic multi-drawing (CLSMD) for the formation of large NF sheets with highly uniform thickness [31,32] and the continuous CLSMD apparatus has also been equipped with an 80-zigzag orifice plate for the formation of long and broad PET NF sheets [33]. Unfortunately, even non-woven NF fabrics have so far been restricted in their application due to their poor mechanical properties, whichever fabrication technique is employed. If the mechanical properties of NFs can be improved, then their application will expand in various fields. This can be achieved by preparing NF yarns, but this is difficult when the mechanical properties of the NF itself are poor. Attempts to produce NF yarns from electrospun NFs have been reported by many researchers. In the early stages, the production of aligned NFs was attempted using a high-speed rotating drum collector and a tapered wheel, and short NF yarns were obtained with such setups [34–36]; however, continuous NF yarns were not successfully produced. Recently, electrospinning setups have been improved by a variety of new techniques and devices, and have been developed to produce continuous NF yarns. As a result, twisted NF yarns with improved mechanical properties have been prepared by twisting an NF web with a rotating disk collector and a rotating funnel [37,38]. We also proposed a technique for the production of continuous PLLA NF multifilaments using a CLSD setup with an interlacer to intermingle NFs and a take-up roll to wind the NF yarn [39]. In these PLLA-NF multifilaments. The low frictional forces among NFs could be lower than twisted fibers because only NFs near the surface of the yarn were confounded by air jet and NFs near the center of NF multifilament cannot did. So
Corresponding author. E-mail address: [email protected] (A. Suzuki).
https://doi.org/10.1016/j.eurpolymj.2018.11.028 Received 21 September 2018; Received in revised form 12 November 2018; Accepted 16 November 2018 Available online 19 November 2018 0014-3057/ © 2018 Elsevier Ltd. All rights reserved.
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Fig. 2. Schematic diagram of the newly designed CLSMD apparatus equipped with a twister-winder for the continuous preparation of TNFY.
measurement. The DSC instrument was calibrated using indium as a standard. The degree of crystallinity (Xc) was determined from the heat of fusion (ΔHm) and the enthalpy of cold crystallization (ΔHcc) as follows:
Fig. 1. WAXD pattern of the original PLLA fiber.
far, multifilaments produced by intermingling have lacked sufficient frictional forces among the constituent NFs. To increase the frictional forces among the NFs and thus improve both the strength and flexibility of the NF yarn, PLLA-NF was twisted using a CLSMD setup equipped with a newly designed and manufactured twister-winder. In this study, continuous twisted nanofiber yarns (TNFYs) with various twist numbers were prepared using a CLSMD setup equipped with a twister-winder, and their superstructure and mechanical properties were characterized before and after annealing at constant length.
Xc =
was used as the heat of fusion for the crystalline where −93 J g phase of PLLA [40]. The mechanical properties of TNFY were determined using a tensile testing machine (Autograph, Shimadzu Co., Japan) with a gauge length of 5 cm and an elongation rate of 10 mm min−1. The average of 10 measurements was taken as an experimental result. Fig. 2 shows a schematic diagram of the CLSMD apparatus newly designed and manufactured for the continuous preparation of TNFY, including a photograph of the winding TNFY. The apparatus consists of a fiber supply spool, nip rolls to supply the original fiber to the injection orifice at a constant speed, a continuous-wave CO2 laser with an output wavelength of 10.6 µm and a maximum power of 40 W, a power meter, a vacuum pump, a vacuum chamber with Zn-Se windows, an orifice plate with 25 fiber injection orifices, a net conveyor to collect the NFs, and a twister-winder that can individually set a winding speed and a twisting speed. The pressure of the vacuum chamber was reduced using a vacuum pump through a linear exhaust port installed below the net conveyor. The laser emitter was placed on a movable platen, which consisted of a micro-alignment stage and a laboratory jack that could be moved parallel to the X- and Z-axes about the laser irradiation point on the fiber, allowing fine adjustments to be made. The vacuum chamber and the orifice plate were coated with black-colored valumite to prevent laser scattering. Continuous PLLA-TNFYs with various twist numbers were prepared by varying the twisting speed and winding speed using the CLSMD setup equipped with the twister-winder. In this study, 20 PLLA fibers injected into the orifices were multi-drawn by irradiating CO2 laser, and NFs collected on the net conveyor were wound while twisting by the twister-winder. Morphologies, superstructure, and mechanical properties of TNFYs obtained were characterized.
The original fiber used in this study was a commercial-grade drawn PLLA fiber with a diameter of 41.6 μm and a 39.6% degree of crystallinity. The original PLLA fiber had a high degree of crystal orientation, as shown by the wide-angle X-ray diffraction (WAXD) pattern in Fig. 1. The equatorial spot reflection in the WAXD pattern is attributable to the (2 0 0)/(1 1 0) reflection from the highly oriented crystallites of the αform, and the degree of crystal orientation estimated using Eq. (1) was 0.960. The morphology of the TNFY was determined using scanning electron microscopy (SEM; JCM-5700, Jeol, Japan). An SEM micrograph of the NF was observed at an accelerating voltage of 10 kV. Prior to observations, the samples were coated with platinum using a sputter coater. The average diameter and the diameter distribution for the NF were obtained using an imaging analyzer. A WAXD pattern of the TNFY was obtained using an imaging-plate (IP) film and an IP detector (R-AXIS DS3C, Rigaku Co., Japan). The IP film was attached to an X-ray generator (Rigaku Co., Japan) operated at 40 kV and 200 mA. The radiation used was Ni-filtered Cu Kα radiation. The sample-to-film distance was 40 mm. The fiber was exposed for 60 min to the X-ray beam from a pinhole collimator with a diameter of 0.4 mm. The degree of crystal orientation (π) was estimated from the half-width (H) of the meridian reflection peak in the WAXD pattern measured with the imaging-plate using data analysis software. The π value is given by:
180 − H × 100 180
(2)
−1
2. Experimental
π=
ΔHm + ΔHcc × 100 −93
3. Results and discussion The morphologies of nanofibers prepared by the CLSMD technique are dependent on the conditions employed, such as the laser power, chamber pressure, fiber supply speed, and laser irradiation position, as reported in the previous papers [24–33], and the laser power and chamber pressure have a significant influence on the fiber diameter and superstructure. In this study, PLLA-TNFYs were prepared by variation of only the twisting conditions, while the laser irradiation conditions
(1)
Differential scanning calorimetry (DSC; Thermo Plus EV02 DSC 8231, Rigaku Co., Japan) measurements were performed within the temperature range of 25–200 °C using a heating rate of 10 °C min−1. All DSC experiments were performed under a nitrogen purge. An approximately 2 mg sample of NFMF was sealed in an aluminum pan for 146
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Fig. 3. Photographs of collected NFs separating from the net conveyor at various winding speeds (SW) and twisting speeds (ST).
Fig. 4. SEM micrographs (50×) of TNFYs obtained at various winding speeds (SW) and twisting speeds (ST), along with the twist number (nT = ST/SW).
3.1. Morphologies of TNFYs prepared by variation of the winding speed and twisting speed
were held constant, and the relationship between the morphology and the twisting conditions was examined closely. Furthermore, the obtained PLLA-TNFYs were annealed to improve their mechanical properties, and their superstructure and mechanical properties were investigated.
Fig. 3 shows photographs of the separation of collected NFs from the net conveyor at various winding speeds (SW) and twisting speeds (ST). All TNFYs were taken up with a right-handed twist. The laser power, chamber pressure, fiber supply speed, and conveyor speed (Scon) were held at 40 W, 30 kPa, 0.20 m min−1, and 0.25 m min−1 respectively,
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during this series of experiments. Because the TNFY was collected on the net conveyor at Scon = 0.25 m min−1, the draw ratio (λ) can be estimated using the following relationship:
λ = S W /Scon
(3)
TNFYs obtained by varying SW have λ values in the range of 1.14–1.76 times. During winding-twisting at these low draw ratios, the NF itself only arranged into the twist direction and was not post-drawn during this process. The position of NF separation from the net conveyor depends on SW and ST, and tends to shift toward the laser emitter (the right side of the photographs) as SW and ST increase. However, under winding conditions of SW = 0.286 m min−1 and ST = 174 T min−1, the separation position moved to the lower side of the net conveyor, and a continuous TNFY could not be obtained because the NFs could not be stably separated from the net conveyor. The starting point of the twisting approached the separation position during NF winding at lower Sw and higher ST. The starting point of the twisting and the separation position were coincident for winding at ST = 300 and 365 T min−1 at Sw = 0.286 m min−1 and in winding at ST = 365 T min−1 and Sw = 0.330 m min−1. Twisting at higher ST contributed to the separation of NFs from the net conveyor under these conditions. Fig. 4 shows SEM micrographs (×50) of TNFYs obtained at various SWs and STs, including the twist number (nT) estimated by
nT = S T /S w
(4)
The estimated nT at each SW increases as ST increases, and the nT at each ST decreases as SW increases. TNFYs with a wide nT range of 395–1276 T m−1 can be obtained by varying SW and ST, but TNFY was not produced when the twisting was carried out at the slowest ST and SW. Fig. 5 shows SEM micrographs (×1000) of TNFYs along with corresponding average fiber diameters (dav), except at SW = 0.286 m min−1 and ST = 174 T min−1. The SEM micrographs show that the NFs had smooth surfaces without surface roughening by laser ablation and there was no droplet formation.
Fig. 6. The twist speed (ST) dependence of the twisted nanofiber yarn diameter (dNFY), twist angle (ϕT), and the bulk density (ρB) for TNFYs obtained at various ■:SW = 0.330 m min−1, twisting conditions; □:SW = 0.286 m min−1, −1 −1 ○:SW = 0.385 m min , ●:SW = 0.440 m min .
Fig. 5. SEM micrographs (1000×) of TNFYs along with the average fiber diameter (dav). 148
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the crystallites because the highly-oriented crystalline nuclei cannot further crystallize due to quenching, even though highly oriented crystalline nuclei were formed in the supersonic jet. PLLA crystallizes in two polymorphic forms, α-form (orthorhombic) and β-form (trigonal) [41–43]. The melting peak temperature (Tm) of the α-form is approximately 10 °C higher than that of the β-form, and the α-form has a 10/7 helix. The original fiber had double melting peaks, but the lower Tm of the original PLLA fiber was not due to the βform because the temperature difference between the low and high Tm was only 6.1 °C. It is instead attributable to an α′-form, so that both αand α′-forms coexist in the original PLLA fiber. The crystal structures of the α- and α′-forms are quite similar, except for the chain conformation and chain packing mode [44]. All TNFYs had only a single sharp melting peak at approximately 164 °C, and their Tms were almost the same as the higher Tm of the original fiber. These melting peaks can be mostly ascribed to α-form crystals that crystallized during the DSC measurement, because all TNFYs had exothermic peaks due to cold crystallization at about 73 °C. These DSC curves suggest that the molecular chains were highly oriented by CLSMD and crystallized easily to act as an α-form nucleating agent, and that α-form nuclei with a higher degree of packing grew during the DSC measurements. Table 1 lists Tm and Xc values estimated from Eq. (2) for TNFYs obtained under various twisting conditions. The Xc values of all TNFYs reached about 39%, which is about equal to that (39.6%) of the original fiber with highly oriented crystallites. Xc depends only upon the laser irradiation conditions and is independent of the twisting conditions. Tm was also independent of the twisting conditions and was held constant at about 164 °C. The observed Tm and Xc values indicate that the superstructure of the TNFYs did not change even after varying the twisting conditions because the superstructure of the NF obtained by CO2 laser supersonic drawing depends only on the laser irradiation conditions. These results suggest that the NFs obtained by CLSMD formed nuclei and/or crystallized only during the super-drawing in the supersonic air jet. The twist of the NF does not affect the formation of nuclei or the crystallization.
The dav values of all TNFYs were in the range of 370–400 nm and were independent of twisting conditions because the dav of the NF depended only upon the laser-irradiation conditions. Fig. 6 shows the ST dependences of the twisted nanofiber yarn diameter (dNFY), twist angle (ϕT), and bulk density (ρB) for TNFYs obtained under various twisting conditions. ρB was estimated from the weight per unit length. The dNFY of TNFY wound at each SW decreased as ST increased, and the dNFY of TNFY twisted at a faster ST was lower than that at a slower ST. The dNFY value of TNFYs obtained at faster SW and ST values decreased because NFs wound at the higher SW were arranged in the winding direction and were strongly tightened in the lateral direction by twisting at the faster ST. ϕT of the TNFY at each SW increased linearly as ST increased, and the ϕT values of TNFYs obtained at the same ST value tended to decrease with increasing SW. The ϕT of TNFYs wound at faster SW were smaller than those at slower SW because the NFs were arranged in the winding direction by winding at the faster SW despite twisting at the same ST. ρB at each SW increased in proportion to ST and tended to increase because the NFs were densely packed by winding at the faster SW. The degree of NF packing depends on ST, and NFs are packed most closely by twisting at the fastest ST. Fig. 7 shows DSC curves for the original fiber and TNFYs twisted at various SW and ST values, except for SW = 0.286 m min−1 and ST = 174 T min−1. The original fiber had a trace with a glass transition temperature (Tg) around 70 °C and double melting peaks at 158.1 and 164.2 °C but did not show cold crystallization (Tcc). All TNFYs had a glass transition around 60 °C, an exotherm at about 73 °C due to cold crystallization, and only one sharp melting peak at approximately 164 °C. The glass transition expresses the degree of amorphous chain movement restriction by crystallites, crystalline nuclei, and entanglements. Tg of the original fiber was 10 °C higher than those of the TNFYs, so the amorphous chains in the original fiber were more strongly restricted by the highly oriented crystallites, as compared to in the TNFY. On the other hand, the lower Tg of the TNFY suggests that the motion of the amorphous chains in the TNFY were not as restricted by
Fig. 7. DSC curves for the original fiber and TNFYs twisted at various winding speeds (SW) and twisting speeds (ST). 149
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TNFYs obtained at ST = 300 T min−1 have twist angles in the range of 30–44°, as shown in Fig. 6. The relationship between Young’s modulus and ϕT in the twisted nanofiber yarns has been reported previously, and in that work, the twist angle giving the maximum Young’s modulus was 35° [45]. At higher ST, the hardness of the TNFY decreased because the degree of crimp in the TNFY increases with increasing twist number (nT = ST/SW). This decrease in the TNFY hardness leads to a decrease in Young’s modulus. The tensile strength of the TNFY obtained at each SW increased with increasing twisting speed, and the tensile strength at each ST increased with increasing SW. The increment of both ST and SW led to an increase in the frictional force among the NFs, directly increasing the tensile strength. The elongation at break increased with increasing twisting speed because the NF length per unit length of TNFY increased and the degree of crimp increased with the increased twisting. The mechanical properties of the twisted yarns depended upon the state of twisting, such as the twist number and the frictional force among NFs, and the effects of these two factors on the mechanical properties of the twisted yarns are very complicated.
Table 1 Melting temperature (Tm) and degree of crystallinity (XC) for TNFYs obtained at various winding speeds (SW) and twisting speeds (ST). Tm/°C Sw/m min−1 ST/T min−1
174 230 300 365
XC/%
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
– 164.2 164.1 163.2
164.4 163.6 163.6 163.4
163.4 164.4 163.8 163.7
163.2 163.4 164.5 163.4
– 39.4 40.0 38.8
39.6 39.4 39.2 39.1
38.9 38.5 37.8 38.9
38.7 38.4 39.3 39.5
3.3. PLLA-TNFY annealed at constant length Annealing is effective at improving the mechanical properties of fibers because crystallization proceeds during annealing under optimum conditions. We reported in a previous study that the mechanical properties of PLLA-NFMF were improved by annealing at constant length [39]. In this study, to improve the mechanical properties of TNFY, yarns obtained by varying ST and SW were annealed at constant length to prevent their untwisting. Fig. 9 shows SEM micrographs (×50 and ×1000) of TNFYs twisted at various SW and ST values and annealed at 110 °C for 20 min at a constant length. As in the previous figures, no SEM micrograph is shown in Fig. 10 for SW = 0.286 m min−1 and ST = 174 T min−1. Table 2 lists the average fiber diameter (dav), yarn diameter (dNFY), twist angle (ϕT), and bulk density (ρB) of the resulting annealed TNFYs. The TNFYs were annealed without welding among the NFs, as shown in Fig. 9. dav of the NFs was relatively unchanged by the annealing, but the dNFY and ϕT values of all the annealed TNFYs decreased after annealing, and ρB increased despite annealing at a constant length without stressing. A 50 cm length of TNFY was sheathed in an aluminum pipe 95 mm long with an outer diameter of 85 mm and wall thickness of 2 mm without stressing. The TNFY after annealing was elongated to about 52 cm without causing thermal shrinkage, even though no external stress was applied during the annealing. The decrease of dNFY and ϕT values suggest that the entire TNFY extended, and that the NFs became more aligned and more closely packed due to the spontaneous elongation of NFs during annealing. Fig. 10 shows DSC curves for the original fiber and the annealed TNFYs, and Table 3 lists Tm and Xc values estimated from Eq. (2) for the annealed TNFYs. All of the annealed TNFYs had only one sharp melting peak at approximately 164 °C but did not show a glass transition or cold crystallization. These melting peaks can be ascribed to α-form crystals that crystallize during the annealing. The Xc values of all the annealed TNFYs reached about 43%, approximately 4% higher than those (39%) of the as-twisted NFYs and the original fiber with highly oriented crystallites. The annealed TNFYs have higher Xc values than the original fiber despite constant-length annealing without stressing. The highly oriented molecular chains produced during CLSMD act as α-form nuclei, and these highly oriented nuclei crystallize easily along the fiber axis during annealing, so that the annealed TNFY has a higher Xc than the original fiber. This crystallization along the fiber axis induced spontaneous elongation of NFs. The presence of these highly oriented crystalline nuclei in the
Fig. 8. ST dependence of Young’s modulus, tensile strength, and elongation at break for TNFYs twisted at various winding speeds (SW) and twisting speeds (ST).
3.2. Mechanical properties of PLLA-TNFY The mechanical properties of TNFY depend only on the twist conditions because the laser irradiation conditions were kept constant and the superstructure, such as Xc and Tm, of all NFs in the twisted yarns was the same, as discussed above. Fig. 8 shows the ST dependence of Young’s modulus, tensile strength, and elongation at break for TNFYs twisted under varying SW and ST values. The Young’s modulus values of TNFYs twisted at SW = 0.330 and 0.440 m min−1 initially increased with increasing ST, and then decreased after the maximum Young’s modulus at each SW was reached. The Young’s modulus at each Sw except for SW = 0.286 m min−1 was obtained in TNFYs twisted at ST = 300 T min−1. The TNFY had the maximum Young’s modulus at an optimum obliquity because the high twist reduced the contribution of the fiber strength to the yarn hardness due to fiber obliquity in the yarn. 150
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Fig. 9. SEM micrographs (50× and 1000×) of TNFYs twisted at various winding speeds (SW) and twisting speeds (ST), annealed at 110 °C for 20 min at a constant length.
Fig. 10. DSC curves for the original fiber and the annealed TNFYs. 151
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Table 2 Average fiber diameter (dav), yarn diameter (dNFY), twist angle (ϕT), and balk density (ρB) for TNFYs, which were obtained at various winding speeds (SW) and twisting speeds (ST), annealed at annealing temperature of 110 °C for 20 min at a constant length. dav/μm Sw/m min−1 ST/T min−1
174 230 300 365
ρB/g cm−3
φT/°
dNFY/μm
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
– 0.384 0.381 0.395
0.381 0.388 0.384 0.384
0.380 0.387 0.383 0.393
0.388 0.383 0.379 0.392
– 368.5 339.5 311.3
382.5 288.8 288.6 267.5
337.0 278.6 260.3 257.2
275.4 264.7 237.4 199.8
– 26.6 36.9 37.6
21.9 26.4 33.6 34.9
21.7 22.4 28.0 36.0
19.0 20.3 25.1 32.2
– 0.273 0.335 0.419
0.197 0.361 0.378 0.472
0.221 0.368 0.431 0.506
0.299 0.349 0.439 0.709
Table 3 Melting temperature (Tm) and degree of crystallinity (XC) for TNFYs, which were obtained at various winding speeds (SW) and twisting speeds (ST), annealed at annealing temperature of 110 °C for 20 min at a constant length. Tm/°C Sw/m min−1 ST/T min−1
174 230 300 365
XC/%
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
– 163.8 163.1 163.6
163.2 164.3 163.7 163.7
162.3 163.2 163.7 163.3
163.9 164.1 163.0 163.1
– 44.0 43.2 43.0
41.3 44.2 433 45.2
43.3 42.6 43.5 42.8
43.5 42.1 42.2 42.6
direction of the fiber axis of the prepared CLSD NFs could not be confirmed because the nuclei are extremely difficult to detect by WAXD. In previous studies [24,39], the formation of highly oriented crystalline nuclei during CLSD was confirmed by collecting NFs in a bundle under the fiber supplying orifice and then annealing the bundled NFs at constant length to reveal the presence of the oriented crystallites. WAXD measurements were carried out for the annealed bundled-NF. The resulting WAXD patterns showed an equatorial arc reflection due to the oriented crystallites, and the reflection was attributed to the reflection of the α-form crystal. The degree of crystal orientation (π) was estimated from the half-width (H) of the meridian of the reflection peak to be 0.88, indicating that the annealed bundled-NF contained highly oriented crystallites. The WAXD measurements experimentally verified that the CLSD-NF contained crystalline nuclei that were highly oriented along the fiber axis, and that the flow induced the formation of these highly oriented crystalline nuclei during CLSD. Fig. 11 shows WAXD patterns from annealed TNFYs obtained using various twisting conditions, along with the π value of the (2 0 0)/(1 1 0) lattice, which was estimated using Eq. (1). The WAXD pattern of the annealed TNFY shows the equatorial arc reflection due to the oriented crystallites, and its reflection is attributable to the (2 0 0)/(1 1 0) reflection of the crystal. The equatorial arc reflection converges at the equator with increasing SW and decreasing ST. The π value expresses the obliquity of the NFs to the TNFY axis because the crystallites present in the NF are highly oriented along the direction of the NF axis, as mentioned above. The π value at each Sw showed a decreasing trend with increasing ST, and that at each ST tended to increase as SW increased. The NFs were strongly twisted at the highest ST and the lowest SW, and the obtained TNFY had the maximum nT = 1276 T m−1 and the lowest ϕT value. The yarn obtained by annealing the TNFY with the maximum nT and ϕT values had the lowest π value. Fig. 12 shows the relationship between ϕT and π for the annealed TNFYs. All of the experimental data points except SW = 0.286 m min−1 are approximately on a straight line, and there is a linear relationship between π and ϕT. When ϕT is extrapolated to 0, π becomes 1, representing perfect orientation of the crystallites, indicating that the NFs were firmly twisted.
Fig. 11. WAXD patterns of yarns obtained by annealing TNFYs obtained at various twisting conditions, along with the degree of crystal orientation (π) of the (2 0 0)/(1 1 0) lattice.
Fig. 12. Relationship between the twist angle (ϕT) and the degree of crystal orientation (π) for annealed TNFYs.
3.4. Mechanical properties of annealed TNFY Fig. 13 shows the ST dependence of the Young’s modulus, tensile strength, and elongation at fracture for annealed TNFYs. The Young’s moduli of all TNFYs was increased by annealing, and were approximately 3–5 times those of the as-twisted TNFYs. The maximum Young’s modulus was obtained by annealing a TNFY twisted at SW = 0.44 m min−1 and ST = 230 T min−1, and was 16.1 cN dtex−1, 5.5 times the Young’s modulus before annealing. The tensile strength of the annealed TNFYs at each SW were 152
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winder. The mechanical properties of TNFYs were changed by varying the twist conditions, and the Young’s modulus had a maximum value in the twist angle range of 30–44°. The tensile strength increased with increasing frictional forces among the NFs caused by increasing both ST and SW. The mechanical properties of the TNFYs increased after annealing at constant length. The Young’s moduli of all the annealed TNFYs increased to approximately 3–5 times those of the as-twisted TNFYs. The tensile strength of the annealed TNFYs at each SW was relatively independent of ST, and the tensile strengths of the annealed TNFYs increased to approximately 1.3–1.5 times those of the as-twisted TNFYs. This method could be applied to produce TNFYs of various polymers, such as PET, nylon, PGA, biodegradable copolymer, and solvent insoluble polymers, without the use of solvents. The resultant TNFYs are flexible and have sufficient mechanical properties for use as a suture material. PLLA-TNFY such as that produced in this study using an apparatus equipped with a twister-winder will be most useful in medical applications because the production process does not require the use of solvents that could be toxic to the human body. Appendix A. Supplementary material Supplementary data to this article can be found online at https:// doi.org/10.1016/j.eurpolymj.2018.11.028. References [1] Y. Yeo, D. Jeon, C. Kim, S. Choi, K. Cho, Y. Lee, C. Kim, J. Biomed. Mater. Res. 72B (2005) 86–93. [2] K. Lee, S. Givens, D.B. Chase, J.F. Rabolt, Polymer 47 (2006) 8013–8018. [3] X. Zong, K. Kim, D. Fang, S. Ran, B.S. Hsiao, B. Chu, Polymer 43 (2002) 4403–4412. [4] J. Meng, L. Song, J. Meng, H. Kong, G. Zhu, C. Wang, L. Xu, S. Xie, H. Xu, J. Biomed. Mater. Res. 79A (2006) 298–306. [5] Y. You, B.M. Min, S.J. Lee, T.S. Lee, W.H. Park, J. Appl. Polym. Sci. 95 (2005) 193–200. [6] B.S. Kim, D.J. Mooney, J. Biomed. Mater. Res. 41 (1998) 322–332. [7] S.P. Higgins, A.K. Solan, L.E. Niklason, J. Biomed. Mater. 67A (2003) 295–302. [8] J. Gao, L. Niklason, R. Langer, J. Biomed. Mater. Res. 42 (1998) 417–424. [9] D.B. Eugene, A.T. Todd, G.S. David, E.W. Gary, L.B. Gary, J. Biomed. Mater. Res. 71B (2004) 144–152. [10] B. Ding, E. Kimura, T. Sato, S. Fujita, S. Shiratori, Polymer 45 (2004) 1895–1902. [11] P. Gupta, G.L. Wilkes, Polymer 44 (2003) 6353–6359. [12] J. Ayutsede, M. Gandhi, S. Sukigara, M. Micklus, H.E. Chen, F. Ko, Polymer 46 (2005) 1625–1634. [13] H. Fong, Polymer 45 (2004) 2427–2432. [14] J.S. Kim, D.H. Reneker, Polym. Eng. Sci. 38 (1999) 849–854. [15] J.M. Deitzel, J. Kleinmeyer, D. Harris, B.N.C. Tan, Polymer 42 (2001) 261–272. [16] C. Huang, S. Chen, D.H. Reneker, C. Lai, H. Hou, Adv. Mater. 18 (2006) 668–671. [17] A. Pedicini, R.J. Farris, Polymer 44 (2003) 6857–6862. [18] J.S. Varabhas, G.G. Chase, D.H. Reneker, Polymer 49 (2008) 4226–4229. [19] H. Zhou, T.B. Green, Y. Joo, Polymer 47 (2006) 7497–7505. [20] P.D. Dalton, D. Grafahrend, K. Klinkhammer, D. Klee, M. Möller, Polymer 48 (2007) 6823–6833. [21] J. Lyons, C. Li, F. Ko, Polymer 45 (2004) 7597–7603. [22] C.J. Ellison, A. Phatak, D.W. Giles, C.W. Macosko, F.S. Bates, Polymer 48 (2007) 3306–3316. [23] S. Borkar, B. Gu, M. Dirmyer, R. Delicado, A.N. Sen, B.R. Jackson, J.V. Badding, Polymer 47 (2006) 8337–8343. [24] A. Suzuki, K. Aoki, Euro Polym. J. 44 (2008) 2499–2505. [25] A. Suzuki, K. Tanizawa, Poymer 50 (2009) 913–921. [26] A. Suzuki, Y. Yamada, J. Appl. Polym. Sci. 116 (2010) 1913–1919. [27] A. Suzuki, R. Shimizu, Appl. Polym. Sci. 121 (2011) 3078–3084. [28] A. Suzuki, H. Hayashi, eXPRESS Polym. Lett. 7 (6) (2013) 519–527. [29] A. Suzuki, T. Mikuni, T. Hasegawa, J. Appl. Polym. Sci. 131 (2014) 40015. [30] H. Koyama, Y. Watanabe, A. Suzuki, J. Polym. Eng. 37 (2017) 53–60. [31] A. Suzuki, K. Arino, Poymer 51 (2010) 1830–1836. [32] A. Suzuki, K. Arino, Euro Polym. J. 48 (2012) 1169–1176. [33] A. Suzuki, K. Hosoi, K. Miyagi, Poymer 60 (2015) 252–259. [34] P.D. Dalton, D. Klee, M. Maller, Polymer 46 (2005) 611–614. [35] A.F. Lotus, E.T. Bender, E.A. Evans, R.D. Ramsier, D.H. Reneker, G.G. Chase, J. Appl. Phys. 103 (2008) 024910–024910 6. [36] L.Q. Liu, M. Eder, I. Burgert, D. Tasis, M. Prato, Wagner H. Daniel, Appl. Phys. Lett. 90 (2007) 083108–083108 3. [37] M. Afifi Amalina, S. Nakano, H. Yamane, Y. Kimura, Macromol. Mater. Eng. 295 (2010) 660–665. [38] U. Ali, Y. Zhou, X. Wang, T. Lin, J. Text. Inst. 103 (2012) 80–88.
Fig. 13. ST dependence of Young’s modulus, tensile strength, and elongation at break for annealed TNFYs; □:SW = 0.286 m min−1, ■:SW = 0.330 m min−1, ○:SW = 0.385 m min−1, ●:SW = 0.440 m min−1.
relatively independent of ST, and the tensile strengths of the annealed TNFYs increased to approximately 1.3–1.5 times those of the as-twisted TNFYs. The elongation at break for the annealed TNFYs was reduced to half that of the as-twisted TNFYs and increased with increasing ST. Improvement in the mechanical properties by annealing TNFYs can be attributed to the increasing of Xc at each NF and the increment in the friction among NFs.
4. Conclusions We have shown that continuous PLLA-TNFY can be successfully obtained using a CO2-laser supersonic multi-drawing apparatus equipped with a twister-winder. TNFYs with various yarn diameters, twist numbers, twist angles, and bulk densities were produced by varying the twisting speed and the winding speed, and TNFYs with a wide nT range of 395–1276 T m−1 were obtained. The yarn diameter of the TNFY at each winding speed decreased as the twisting speed increased, and the twist angle and bulk density increased with increasing twist speed. The fiber diameter of the individual nanofibers comprising a TNFY depended only upon the laser-irradiation conditions and was independent of twist conditions. The nanofiber diameters of all TNFYs were in the range from 370 to 400 nm and were almost equal because the laser-irradiation conditions were kept constant in the preparation of the TNFYs. In the relationship between the twist angle and the degree of crystal orientation for almost of the annealed TNFYs, the experimental data points (except SW = 0.286 m min−1) fell approximately on a straight line, suggesting that the NFs were firmly twisted by the twister153
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[39] [40] [41] [42]
A. Suzuki, K. Imajo, Poymer 91 (2016) 24–32. B. Kalb, A.J. Pennings, Polymer 21 (1980) 607–612. B. Eling, S. Gogolewski, A.J. Pennings, Polymer 23 (1982) 1587–1593. W. Hoogsteen, A.R. Postema, A.J. Pennings, G.T. Brinke, P. Zugenmaier, Macromolecules 23 (1990) 634–642.
[43] T. Miyata, T. Masuko, Polymer 38 (1997) 4003–4009. [44] D. Dikovsky, G. Marom, C.A. Avila-Orta, R.H. Somani, B.S. Hsiao, Polymer 46 (2005) 3096–3104. [45] H. Yan, L. Liu, Z. Zhang, Mater. Lett. 65 (2011) 2419–2421.
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Poly(L-lactic acid) twisted nanofiber yarn prepared by carbon dioxide laser supersonic multi-drawing
T
⁎
Akihiro Suzuki , Yuta Shimba University of Yamanashi, Interdisciplinary Graduate of School, Takeda-4, Kofu 400-8511, Japan
A B S T R A C T
Poly(L-lactic acid) (PLLA) twisted nanofiber yarn (TNFY) was continuously prepared using a CO2-laser supersonic multi-drawing apparatus equipped with a twisterwinder that offers independent control of winding speed (SW) and twisting speed (ST). TNFYs with various twist numbers (nT), twist angles, and bulk densities were produced by changing SW and ST. TNFYs with a wide nT range of 395–1276 T m−1 were obtained. Furthermore, the mechanical properties of the obtained TNFYs depended upon the degree of twist. The maximum Young’s modulus was 3.06 cN dtex−1 at ST = 300 T min−1 and SW = 0.440 m min−1, and the maximum tensile strength was 0.36 cN dtex−1 at ST = 365 T min−1 and SW = 0.440 m min−1. The Young’s modulus had a maximum value at optimum obliquity because the high twist reduced the contribution of fiber strength to the yarn hardness due to fiber obliquity in the yarn. The tensile strength increased with increasing ST because the frictional forces among the NFs were increased by twisting at higher ST.
1. Introduction Nanofibers (NFs) are used in a variety of applications, such as membranes [1–3], biomedical devices [4], and scaffolds for tissue engineering [5–9]. NFs have been produced by electrospinning [10–21], sea-island-type conjugated melt spinning, and melt blowing [22,23]. We previously proposed a technique for the production of NFs, referred to as a carbon dioxide (CO2) laser supersonic drawing (CLSD). It involves irradiating a fiber with radiation from CO2-laser while drawing it at a supersonic velocity. A supersonic jet was generated by blowing air into a vacuum chamber through the orifice used to inject the fiber into the vacuum chamber. The adiabatic expansion of air across the orifice cools the jet. The fiber is instantly melted by the high-power laser beam that irradiates in the cold supersonic jet. It is then tremendously deformed by the shear force generated by the supersonic flow, and ultradrawn to a draw ratio of the order of 105. CLSD is a novel technique that uses laser irradiation of fibers in a supersonic jet, and can easily prepare NFs of various polymers. CLSD employs only CO2 laser irradiation and does not require additional processes or solvents, unlike electrospinning, which requires a solvent, or sea-island-type conjugated melt spinning that requires the removal of a second component. CLSD has already been applied to prepare NFs of poly(L-lactic acid) (PLLA) [24], poly(ethylene terephthalate) [25], poly (ethylene-2,6-naphthalate) [26], poly(glycolic acid) [27], ethylene tetrafluoroethylene copolymer [28], nylon 66 [29], and poly(phenylene sulfide) [30]. Modification of the CLSD technique has led to the development of
⁎
CO2-laser supersonic multi-drawing (CLSMD) for the formation of large NF sheets with highly uniform thickness [31,32] and the continuous CLSMD apparatus has also been equipped with an 80-zigzag orifice plate for the formation of long and broad PET NF sheets [33]. Unfortunately, even non-woven NF fabrics have so far been restricted in their application due to their poor mechanical properties, whichever fabrication technique is employed. If the mechanical properties of NFs can be improved, then their application will expand in various fields. This can be achieved by preparing NF yarns, but this is difficult when the mechanical properties of the NF itself are poor. Attempts to produce NF yarns from electrospun NFs have been reported by many researchers. In the early stages, the production of aligned NFs was attempted using a high-speed rotating drum collector and a tapered wheel, and short NF yarns were obtained with such setups [34–36]; however, continuous NF yarns were not successfully produced. Recently, electrospinning setups have been improved by a variety of new techniques and devices, and have been developed to produce continuous NF yarns. As a result, twisted NF yarns with improved mechanical properties have been prepared by twisting an NF web with a rotating disk collector and a rotating funnel [37,38]. We also proposed a technique for the production of continuous PLLA NF multifilaments using a CLSD setup with an interlacer to intermingle NFs and a take-up roll to wind the NF yarn [39]. In these PLLA-NF multifilaments. The low frictional forces among NFs could be lower than twisted fibers because only NFs near the surface of the yarn were confounded by air jet and NFs near the center of NF multifilament cannot did. So
Corresponding author. E-mail address: [email protected] (A. Suzuki).
https://doi.org/10.1016/j.eurpolymj.2018.11.028 Received 21 September 2018; Received in revised form 12 November 2018; Accepted 16 November 2018 Available online 19 November 2018 0014-3057/ © 2018 Elsevier Ltd. All rights reserved.
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Fig. 2. Schematic diagram of the newly designed CLSMD apparatus equipped with a twister-winder for the continuous preparation of TNFY.
measurement. The DSC instrument was calibrated using indium as a standard. The degree of crystallinity (Xc) was determined from the heat of fusion (ΔHm) and the enthalpy of cold crystallization (ΔHcc) as follows:
Fig. 1. WAXD pattern of the original PLLA fiber.
far, multifilaments produced by intermingling have lacked sufficient frictional forces among the constituent NFs. To increase the frictional forces among the NFs and thus improve both the strength and flexibility of the NF yarn, PLLA-NF was twisted using a CLSMD setup equipped with a newly designed and manufactured twister-winder. In this study, continuous twisted nanofiber yarns (TNFYs) with various twist numbers were prepared using a CLSMD setup equipped with a twister-winder, and their superstructure and mechanical properties were characterized before and after annealing at constant length.
Xc =
was used as the heat of fusion for the crystalline where −93 J g phase of PLLA [40]. The mechanical properties of TNFY were determined using a tensile testing machine (Autograph, Shimadzu Co., Japan) with a gauge length of 5 cm and an elongation rate of 10 mm min−1. The average of 10 measurements was taken as an experimental result. Fig. 2 shows a schematic diagram of the CLSMD apparatus newly designed and manufactured for the continuous preparation of TNFY, including a photograph of the winding TNFY. The apparatus consists of a fiber supply spool, nip rolls to supply the original fiber to the injection orifice at a constant speed, a continuous-wave CO2 laser with an output wavelength of 10.6 µm and a maximum power of 40 W, a power meter, a vacuum pump, a vacuum chamber with Zn-Se windows, an orifice plate with 25 fiber injection orifices, a net conveyor to collect the NFs, and a twister-winder that can individually set a winding speed and a twisting speed. The pressure of the vacuum chamber was reduced using a vacuum pump through a linear exhaust port installed below the net conveyor. The laser emitter was placed on a movable platen, which consisted of a micro-alignment stage and a laboratory jack that could be moved parallel to the X- and Z-axes about the laser irradiation point on the fiber, allowing fine adjustments to be made. The vacuum chamber and the orifice plate were coated with black-colored valumite to prevent laser scattering. Continuous PLLA-TNFYs with various twist numbers were prepared by varying the twisting speed and winding speed using the CLSMD setup equipped with the twister-winder. In this study, 20 PLLA fibers injected into the orifices were multi-drawn by irradiating CO2 laser, and NFs collected on the net conveyor were wound while twisting by the twister-winder. Morphologies, superstructure, and mechanical properties of TNFYs obtained were characterized.
The original fiber used in this study was a commercial-grade drawn PLLA fiber with a diameter of 41.6 μm and a 39.6% degree of crystallinity. The original PLLA fiber had a high degree of crystal orientation, as shown by the wide-angle X-ray diffraction (WAXD) pattern in Fig. 1. The equatorial spot reflection in the WAXD pattern is attributable to the (2 0 0)/(1 1 0) reflection from the highly oriented crystallites of the αform, and the degree of crystal orientation estimated using Eq. (1) was 0.960. The morphology of the TNFY was determined using scanning electron microscopy (SEM; JCM-5700, Jeol, Japan). An SEM micrograph of the NF was observed at an accelerating voltage of 10 kV. Prior to observations, the samples were coated with platinum using a sputter coater. The average diameter and the diameter distribution for the NF were obtained using an imaging analyzer. A WAXD pattern of the TNFY was obtained using an imaging-plate (IP) film and an IP detector (R-AXIS DS3C, Rigaku Co., Japan). The IP film was attached to an X-ray generator (Rigaku Co., Japan) operated at 40 kV and 200 mA. The radiation used was Ni-filtered Cu Kα radiation. The sample-to-film distance was 40 mm. The fiber was exposed for 60 min to the X-ray beam from a pinhole collimator with a diameter of 0.4 mm. The degree of crystal orientation (π) was estimated from the half-width (H) of the meridian reflection peak in the WAXD pattern measured with the imaging-plate using data analysis software. The π value is given by:
180 − H × 100 180
(2)
−1
2. Experimental
π=
ΔHm + ΔHcc × 100 −93
3. Results and discussion The morphologies of nanofibers prepared by the CLSMD technique are dependent on the conditions employed, such as the laser power, chamber pressure, fiber supply speed, and laser irradiation position, as reported in the previous papers [24–33], and the laser power and chamber pressure have a significant influence on the fiber diameter and superstructure. In this study, PLLA-TNFYs were prepared by variation of only the twisting conditions, while the laser irradiation conditions
(1)
Differential scanning calorimetry (DSC; Thermo Plus EV02 DSC 8231, Rigaku Co., Japan) measurements were performed within the temperature range of 25–200 °C using a heating rate of 10 °C min−1. All DSC experiments were performed under a nitrogen purge. An approximately 2 mg sample of NFMF was sealed in an aluminum pan for 146
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Fig. 3. Photographs of collected NFs separating from the net conveyor at various winding speeds (SW) and twisting speeds (ST).
Fig. 4. SEM micrographs (50×) of TNFYs obtained at various winding speeds (SW) and twisting speeds (ST), along with the twist number (nT = ST/SW).
3.1. Morphologies of TNFYs prepared by variation of the winding speed and twisting speed
were held constant, and the relationship between the morphology and the twisting conditions was examined closely. Furthermore, the obtained PLLA-TNFYs were annealed to improve their mechanical properties, and their superstructure and mechanical properties were investigated.
Fig. 3 shows photographs of the separation of collected NFs from the net conveyor at various winding speeds (SW) and twisting speeds (ST). All TNFYs were taken up with a right-handed twist. The laser power, chamber pressure, fiber supply speed, and conveyor speed (Scon) were held at 40 W, 30 kPa, 0.20 m min−1, and 0.25 m min−1 respectively,
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during this series of experiments. Because the TNFY was collected on the net conveyor at Scon = 0.25 m min−1, the draw ratio (λ) can be estimated using the following relationship:
λ = S W /Scon
(3)
TNFYs obtained by varying SW have λ values in the range of 1.14–1.76 times. During winding-twisting at these low draw ratios, the NF itself only arranged into the twist direction and was not post-drawn during this process. The position of NF separation from the net conveyor depends on SW and ST, and tends to shift toward the laser emitter (the right side of the photographs) as SW and ST increase. However, under winding conditions of SW = 0.286 m min−1 and ST = 174 T min−1, the separation position moved to the lower side of the net conveyor, and a continuous TNFY could not be obtained because the NFs could not be stably separated from the net conveyor. The starting point of the twisting approached the separation position during NF winding at lower Sw and higher ST. The starting point of the twisting and the separation position were coincident for winding at ST = 300 and 365 T min−1 at Sw = 0.286 m min−1 and in winding at ST = 365 T min−1 and Sw = 0.330 m min−1. Twisting at higher ST contributed to the separation of NFs from the net conveyor under these conditions. Fig. 4 shows SEM micrographs (×50) of TNFYs obtained at various SWs and STs, including the twist number (nT) estimated by
nT = S T /S w
(4)
The estimated nT at each SW increases as ST increases, and the nT at each ST decreases as SW increases. TNFYs with a wide nT range of 395–1276 T m−1 can be obtained by varying SW and ST, but TNFY was not produced when the twisting was carried out at the slowest ST and SW. Fig. 5 shows SEM micrographs (×1000) of TNFYs along with corresponding average fiber diameters (dav), except at SW = 0.286 m min−1 and ST = 174 T min−1. The SEM micrographs show that the NFs had smooth surfaces without surface roughening by laser ablation and there was no droplet formation.
Fig. 6. The twist speed (ST) dependence of the twisted nanofiber yarn diameter (dNFY), twist angle (ϕT), and the bulk density (ρB) for TNFYs obtained at various ■:SW = 0.330 m min−1, twisting conditions; □:SW = 0.286 m min−1, −1 −1 ○:SW = 0.385 m min , ●:SW = 0.440 m min .
Fig. 5. SEM micrographs (1000×) of TNFYs along with the average fiber diameter (dav). 148
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the crystallites because the highly-oriented crystalline nuclei cannot further crystallize due to quenching, even though highly oriented crystalline nuclei were formed in the supersonic jet. PLLA crystallizes in two polymorphic forms, α-form (orthorhombic) and β-form (trigonal) [41–43]. The melting peak temperature (Tm) of the α-form is approximately 10 °C higher than that of the β-form, and the α-form has a 10/7 helix. The original fiber had double melting peaks, but the lower Tm of the original PLLA fiber was not due to the βform because the temperature difference between the low and high Tm was only 6.1 °C. It is instead attributable to an α′-form, so that both αand α′-forms coexist in the original PLLA fiber. The crystal structures of the α- and α′-forms are quite similar, except for the chain conformation and chain packing mode [44]. All TNFYs had only a single sharp melting peak at approximately 164 °C, and their Tms were almost the same as the higher Tm of the original fiber. These melting peaks can be mostly ascribed to α-form crystals that crystallized during the DSC measurement, because all TNFYs had exothermic peaks due to cold crystallization at about 73 °C. These DSC curves suggest that the molecular chains were highly oriented by CLSMD and crystallized easily to act as an α-form nucleating agent, and that α-form nuclei with a higher degree of packing grew during the DSC measurements. Table 1 lists Tm and Xc values estimated from Eq. (2) for TNFYs obtained under various twisting conditions. The Xc values of all TNFYs reached about 39%, which is about equal to that (39.6%) of the original fiber with highly oriented crystallites. Xc depends only upon the laser irradiation conditions and is independent of the twisting conditions. Tm was also independent of the twisting conditions and was held constant at about 164 °C. The observed Tm and Xc values indicate that the superstructure of the TNFYs did not change even after varying the twisting conditions because the superstructure of the NF obtained by CO2 laser supersonic drawing depends only on the laser irradiation conditions. These results suggest that the NFs obtained by CLSMD formed nuclei and/or crystallized only during the super-drawing in the supersonic air jet. The twist of the NF does not affect the formation of nuclei or the crystallization.
The dav values of all TNFYs were in the range of 370–400 nm and were independent of twisting conditions because the dav of the NF depended only upon the laser-irradiation conditions. Fig. 6 shows the ST dependences of the twisted nanofiber yarn diameter (dNFY), twist angle (ϕT), and bulk density (ρB) for TNFYs obtained under various twisting conditions. ρB was estimated from the weight per unit length. The dNFY of TNFY wound at each SW decreased as ST increased, and the dNFY of TNFY twisted at a faster ST was lower than that at a slower ST. The dNFY value of TNFYs obtained at faster SW and ST values decreased because NFs wound at the higher SW were arranged in the winding direction and were strongly tightened in the lateral direction by twisting at the faster ST. ϕT of the TNFY at each SW increased linearly as ST increased, and the ϕT values of TNFYs obtained at the same ST value tended to decrease with increasing SW. The ϕT of TNFYs wound at faster SW were smaller than those at slower SW because the NFs were arranged in the winding direction by winding at the faster SW despite twisting at the same ST. ρB at each SW increased in proportion to ST and tended to increase because the NFs were densely packed by winding at the faster SW. The degree of NF packing depends on ST, and NFs are packed most closely by twisting at the fastest ST. Fig. 7 shows DSC curves for the original fiber and TNFYs twisted at various SW and ST values, except for SW = 0.286 m min−1 and ST = 174 T min−1. The original fiber had a trace with a glass transition temperature (Tg) around 70 °C and double melting peaks at 158.1 and 164.2 °C but did not show cold crystallization (Tcc). All TNFYs had a glass transition around 60 °C, an exotherm at about 73 °C due to cold crystallization, and only one sharp melting peak at approximately 164 °C. The glass transition expresses the degree of amorphous chain movement restriction by crystallites, crystalline nuclei, and entanglements. Tg of the original fiber was 10 °C higher than those of the TNFYs, so the amorphous chains in the original fiber were more strongly restricted by the highly oriented crystallites, as compared to in the TNFY. On the other hand, the lower Tg of the TNFY suggests that the motion of the amorphous chains in the TNFY were not as restricted by
Fig. 7. DSC curves for the original fiber and TNFYs twisted at various winding speeds (SW) and twisting speeds (ST). 149
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TNFYs obtained at ST = 300 T min−1 have twist angles in the range of 30–44°, as shown in Fig. 6. The relationship between Young’s modulus and ϕT in the twisted nanofiber yarns has been reported previously, and in that work, the twist angle giving the maximum Young’s modulus was 35° [45]. At higher ST, the hardness of the TNFY decreased because the degree of crimp in the TNFY increases with increasing twist number (nT = ST/SW). This decrease in the TNFY hardness leads to a decrease in Young’s modulus. The tensile strength of the TNFY obtained at each SW increased with increasing twisting speed, and the tensile strength at each ST increased with increasing SW. The increment of both ST and SW led to an increase in the frictional force among the NFs, directly increasing the tensile strength. The elongation at break increased with increasing twisting speed because the NF length per unit length of TNFY increased and the degree of crimp increased with the increased twisting. The mechanical properties of the twisted yarns depended upon the state of twisting, such as the twist number and the frictional force among NFs, and the effects of these two factors on the mechanical properties of the twisted yarns are very complicated.
Table 1 Melting temperature (Tm) and degree of crystallinity (XC) for TNFYs obtained at various winding speeds (SW) and twisting speeds (ST). Tm/°C Sw/m min−1 ST/T min−1
174 230 300 365
XC/%
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
– 164.2 164.1 163.2
164.4 163.6 163.6 163.4
163.4 164.4 163.8 163.7
163.2 163.4 164.5 163.4
– 39.4 40.0 38.8
39.6 39.4 39.2 39.1
38.9 38.5 37.8 38.9
38.7 38.4 39.3 39.5
3.3. PLLA-TNFY annealed at constant length Annealing is effective at improving the mechanical properties of fibers because crystallization proceeds during annealing under optimum conditions. We reported in a previous study that the mechanical properties of PLLA-NFMF were improved by annealing at constant length [39]. In this study, to improve the mechanical properties of TNFY, yarns obtained by varying ST and SW were annealed at constant length to prevent their untwisting. Fig. 9 shows SEM micrographs (×50 and ×1000) of TNFYs twisted at various SW and ST values and annealed at 110 °C for 20 min at a constant length. As in the previous figures, no SEM micrograph is shown in Fig. 10 for SW = 0.286 m min−1 and ST = 174 T min−1. Table 2 lists the average fiber diameter (dav), yarn diameter (dNFY), twist angle (ϕT), and bulk density (ρB) of the resulting annealed TNFYs. The TNFYs were annealed without welding among the NFs, as shown in Fig. 9. dav of the NFs was relatively unchanged by the annealing, but the dNFY and ϕT values of all the annealed TNFYs decreased after annealing, and ρB increased despite annealing at a constant length without stressing. A 50 cm length of TNFY was sheathed in an aluminum pipe 95 mm long with an outer diameter of 85 mm and wall thickness of 2 mm without stressing. The TNFY after annealing was elongated to about 52 cm without causing thermal shrinkage, even though no external stress was applied during the annealing. The decrease of dNFY and ϕT values suggest that the entire TNFY extended, and that the NFs became more aligned and more closely packed due to the spontaneous elongation of NFs during annealing. Fig. 10 shows DSC curves for the original fiber and the annealed TNFYs, and Table 3 lists Tm and Xc values estimated from Eq. (2) for the annealed TNFYs. All of the annealed TNFYs had only one sharp melting peak at approximately 164 °C but did not show a glass transition or cold crystallization. These melting peaks can be ascribed to α-form crystals that crystallize during the annealing. The Xc values of all the annealed TNFYs reached about 43%, approximately 4% higher than those (39%) of the as-twisted NFYs and the original fiber with highly oriented crystallites. The annealed TNFYs have higher Xc values than the original fiber despite constant-length annealing without stressing. The highly oriented molecular chains produced during CLSMD act as α-form nuclei, and these highly oriented nuclei crystallize easily along the fiber axis during annealing, so that the annealed TNFY has a higher Xc than the original fiber. This crystallization along the fiber axis induced spontaneous elongation of NFs. The presence of these highly oriented crystalline nuclei in the
Fig. 8. ST dependence of Young’s modulus, tensile strength, and elongation at break for TNFYs twisted at various winding speeds (SW) and twisting speeds (ST).
3.2. Mechanical properties of PLLA-TNFY The mechanical properties of TNFY depend only on the twist conditions because the laser irradiation conditions were kept constant and the superstructure, such as Xc and Tm, of all NFs in the twisted yarns was the same, as discussed above. Fig. 8 shows the ST dependence of Young’s modulus, tensile strength, and elongation at break for TNFYs twisted under varying SW and ST values. The Young’s modulus values of TNFYs twisted at SW = 0.330 and 0.440 m min−1 initially increased with increasing ST, and then decreased after the maximum Young’s modulus at each SW was reached. The Young’s modulus at each Sw except for SW = 0.286 m min−1 was obtained in TNFYs twisted at ST = 300 T min−1. The TNFY had the maximum Young’s modulus at an optimum obliquity because the high twist reduced the contribution of the fiber strength to the yarn hardness due to fiber obliquity in the yarn. 150
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Fig. 9. SEM micrographs (50× and 1000×) of TNFYs twisted at various winding speeds (SW) and twisting speeds (ST), annealed at 110 °C for 20 min at a constant length.
Fig. 10. DSC curves for the original fiber and the annealed TNFYs. 151
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Table 2 Average fiber diameter (dav), yarn diameter (dNFY), twist angle (ϕT), and balk density (ρB) for TNFYs, which were obtained at various winding speeds (SW) and twisting speeds (ST), annealed at annealing temperature of 110 °C for 20 min at a constant length. dav/μm Sw/m min−1 ST/T min−1
174 230 300 365
ρB/g cm−3
φT/°
dNFY/μm
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
– 0.384 0.381 0.395
0.381 0.388 0.384 0.384
0.380 0.387 0.383 0.393
0.388 0.383 0.379 0.392
– 368.5 339.5 311.3
382.5 288.8 288.6 267.5
337.0 278.6 260.3 257.2
275.4 264.7 237.4 199.8
– 26.6 36.9 37.6
21.9 26.4 33.6 34.9
21.7 22.4 28.0 36.0
19.0 20.3 25.1 32.2
– 0.273 0.335 0.419
0.197 0.361 0.378 0.472
0.221 0.368 0.431 0.506
0.299 0.349 0.439 0.709
Table 3 Melting temperature (Tm) and degree of crystallinity (XC) for TNFYs, which were obtained at various winding speeds (SW) and twisting speeds (ST), annealed at annealing temperature of 110 °C for 20 min at a constant length. Tm/°C Sw/m min−1 ST/T min−1
174 230 300 365
XC/%
0.286
0.330
0.385
0.440
0.286
0.330
0.385
0.440
– 163.8 163.1 163.6
163.2 164.3 163.7 163.7
162.3 163.2 163.7 163.3
163.9 164.1 163.0 163.1
– 44.0 43.2 43.0
41.3 44.2 433 45.2
43.3 42.6 43.5 42.8
43.5 42.1 42.2 42.6
direction of the fiber axis of the prepared CLSD NFs could not be confirmed because the nuclei are extremely difficult to detect by WAXD. In previous studies [24,39], the formation of highly oriented crystalline nuclei during CLSD was confirmed by collecting NFs in a bundle under the fiber supplying orifice and then annealing the bundled NFs at constant length to reveal the presence of the oriented crystallites. WAXD measurements were carried out for the annealed bundled-NF. The resulting WAXD patterns showed an equatorial arc reflection due to the oriented crystallites, and the reflection was attributed to the reflection of the α-form crystal. The degree of crystal orientation (π) was estimated from the half-width (H) of the meridian of the reflection peak to be 0.88, indicating that the annealed bundled-NF contained highly oriented crystallites. The WAXD measurements experimentally verified that the CLSD-NF contained crystalline nuclei that were highly oriented along the fiber axis, and that the flow induced the formation of these highly oriented crystalline nuclei during CLSD. Fig. 11 shows WAXD patterns from annealed TNFYs obtained using various twisting conditions, along with the π value of the (2 0 0)/(1 1 0) lattice, which was estimated using Eq. (1). The WAXD pattern of the annealed TNFY shows the equatorial arc reflection due to the oriented crystallites, and its reflection is attributable to the (2 0 0)/(1 1 0) reflection of the crystal. The equatorial arc reflection converges at the equator with increasing SW and decreasing ST. The π value expresses the obliquity of the NFs to the TNFY axis because the crystallites present in the NF are highly oriented along the direction of the NF axis, as mentioned above. The π value at each Sw showed a decreasing trend with increasing ST, and that at each ST tended to increase as SW increased. The NFs were strongly twisted at the highest ST and the lowest SW, and the obtained TNFY had the maximum nT = 1276 T m−1 and the lowest ϕT value. The yarn obtained by annealing the TNFY with the maximum nT and ϕT values had the lowest π value. Fig. 12 shows the relationship between ϕT and π for the annealed TNFYs. All of the experimental data points except SW = 0.286 m min−1 are approximately on a straight line, and there is a linear relationship between π and ϕT. When ϕT is extrapolated to 0, π becomes 1, representing perfect orientation of the crystallites, indicating that the NFs were firmly twisted.
Fig. 11. WAXD patterns of yarns obtained by annealing TNFYs obtained at various twisting conditions, along with the degree of crystal orientation (π) of the (2 0 0)/(1 1 0) lattice.
Fig. 12. Relationship between the twist angle (ϕT) and the degree of crystal orientation (π) for annealed TNFYs.
3.4. Mechanical properties of annealed TNFY Fig. 13 shows the ST dependence of the Young’s modulus, tensile strength, and elongation at fracture for annealed TNFYs. The Young’s moduli of all TNFYs was increased by annealing, and were approximately 3–5 times those of the as-twisted TNFYs. The maximum Young’s modulus was obtained by annealing a TNFY twisted at SW = 0.44 m min−1 and ST = 230 T min−1, and was 16.1 cN dtex−1, 5.5 times the Young’s modulus before annealing. The tensile strength of the annealed TNFYs at each SW were 152
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winder. The mechanical properties of TNFYs were changed by varying the twist conditions, and the Young’s modulus had a maximum value in the twist angle range of 30–44°. The tensile strength increased with increasing frictional forces among the NFs caused by increasing both ST and SW. The mechanical properties of the TNFYs increased after annealing at constant length. The Young’s moduli of all the annealed TNFYs increased to approximately 3–5 times those of the as-twisted TNFYs. The tensile strength of the annealed TNFYs at each SW was relatively independent of ST, and the tensile strengths of the annealed TNFYs increased to approximately 1.3–1.5 times those of the as-twisted TNFYs. This method could be applied to produce TNFYs of various polymers, such as PET, nylon, PGA, biodegradable copolymer, and solvent insoluble polymers, without the use of solvents. The resultant TNFYs are flexible and have sufficient mechanical properties for use as a suture material. PLLA-TNFY such as that produced in this study using an apparatus equipped with a twister-winder will be most useful in medical applications because the production process does not require the use of solvents that could be toxic to the human body. Appendix A. Supplementary material Supplementary data to this article can be found online at https:// doi.org/10.1016/j.eurpolymj.2018.11.028. References [1] Y. Yeo, D. Jeon, C. Kim, S. Choi, K. Cho, Y. Lee, C. Kim, J. Biomed. Mater. Res. 72B (2005) 86–93. [2] K. Lee, S. Givens, D.B. Chase, J.F. Rabolt, Polymer 47 (2006) 8013–8018. [3] X. Zong, K. Kim, D. Fang, S. Ran, B.S. Hsiao, B. Chu, Polymer 43 (2002) 4403–4412. [4] J. Meng, L. Song, J. Meng, H. Kong, G. Zhu, C. Wang, L. Xu, S. Xie, H. Xu, J. Biomed. Mater. Res. 79A (2006) 298–306. [5] Y. You, B.M. Min, S.J. Lee, T.S. Lee, W.H. Park, J. Appl. Polym. Sci. 95 (2005) 193–200. [6] B.S. Kim, D.J. Mooney, J. Biomed. Mater. Res. 41 (1998) 322–332. [7] S.P. Higgins, A.K. Solan, L.E. Niklason, J. Biomed. Mater. 67A (2003) 295–302. [8] J. Gao, L. Niklason, R. Langer, J. Biomed. Mater. Res. 42 (1998) 417–424. [9] D.B. Eugene, A.T. Todd, G.S. David, E.W. Gary, L.B. Gary, J. Biomed. Mater. Res. 71B (2004) 144–152. [10] B. Ding, E. Kimura, T. Sato, S. Fujita, S. Shiratori, Polymer 45 (2004) 1895–1902. [11] P. Gupta, G.L. Wilkes, Polymer 44 (2003) 6353–6359. [12] J. Ayutsede, M. Gandhi, S. Sukigara, M. Micklus, H.E. Chen, F. Ko, Polymer 46 (2005) 1625–1634. [13] H. Fong, Polymer 45 (2004) 2427–2432. [14] J.S. Kim, D.H. Reneker, Polym. Eng. Sci. 38 (1999) 849–854. [15] J.M. Deitzel, J. Kleinmeyer, D. Harris, B.N.C. Tan, Polymer 42 (2001) 261–272. [16] C. Huang, S. Chen, D.H. Reneker, C. Lai, H. Hou, Adv. Mater. 18 (2006) 668–671. [17] A. Pedicini, R.J. Farris, Polymer 44 (2003) 6857–6862. [18] J.S. Varabhas, G.G. Chase, D.H. Reneker, Polymer 49 (2008) 4226–4229. [19] H. Zhou, T.B. Green, Y. Joo, Polymer 47 (2006) 7497–7505. [20] P.D. Dalton, D. Grafahrend, K. Klinkhammer, D. Klee, M. Möller, Polymer 48 (2007) 6823–6833. [21] J. Lyons, C. Li, F. Ko, Polymer 45 (2004) 7597–7603. [22] C.J. Ellison, A. Phatak, D.W. Giles, C.W. Macosko, F.S. Bates, Polymer 48 (2007) 3306–3316. [23] S. Borkar, B. Gu, M. Dirmyer, R. Delicado, A.N. Sen, B.R. Jackson, J.V. Badding, Polymer 47 (2006) 8337–8343. [24] A. Suzuki, K. Aoki, Euro Polym. J. 44 (2008) 2499–2505. [25] A. Suzuki, K. Tanizawa, Poymer 50 (2009) 913–921. [26] A. Suzuki, Y. Yamada, J. Appl. Polym. Sci. 116 (2010) 1913–1919. [27] A. Suzuki, R. Shimizu, Appl. Polym. Sci. 121 (2011) 3078–3084. [28] A. Suzuki, H. Hayashi, eXPRESS Polym. Lett. 7 (6) (2013) 519–527. [29] A. Suzuki, T. Mikuni, T. Hasegawa, J. Appl. Polym. Sci. 131 (2014) 40015. [30] H. Koyama, Y. Watanabe, A. Suzuki, J. Polym. Eng. 37 (2017) 53–60. [31] A. Suzuki, K. Arino, Poymer 51 (2010) 1830–1836. [32] A. Suzuki, K. Arino, Euro Polym. J. 48 (2012) 1169–1176. [33] A. Suzuki, K. Hosoi, K. Miyagi, Poymer 60 (2015) 252–259. [34] P.D. Dalton, D. Klee, M. Maller, Polymer 46 (2005) 611–614. [35] A.F. Lotus, E.T. Bender, E.A. Evans, R.D. Ramsier, D.H. Reneker, G.G. Chase, J. Appl. Phys. 103 (2008) 024910–024910 6. [36] L.Q. Liu, M. Eder, I. Burgert, D. Tasis, M. Prato, Wagner H. Daniel, Appl. Phys. Lett. 90 (2007) 083108–083108 3. [37] M. Afifi Amalina, S. Nakano, H. Yamane, Y. Kimura, Macromol. Mater. Eng. 295 (2010) 660–665. [38] U. Ali, Y. Zhou, X. Wang, T. Lin, J. Text. Inst. 103 (2012) 80–88.
Fig. 13. ST dependence of Young’s modulus, tensile strength, and elongation at break for annealed TNFYs; □:SW = 0.286 m min−1, ■:SW = 0.330 m min−1, ○:SW = 0.385 m min−1, ●:SW = 0.440 m min−1.
relatively independent of ST, and the tensile strengths of the annealed TNFYs increased to approximately 1.3–1.5 times those of the as-twisted TNFYs. The elongation at break for the annealed TNFYs was reduced to half that of the as-twisted TNFYs and increased with increasing ST. Improvement in the mechanical properties by annealing TNFYs can be attributed to the increasing of Xc at each NF and the increment in the friction among NFs.
4. Conclusions We have shown that continuous PLLA-TNFY can be successfully obtained using a CO2-laser supersonic multi-drawing apparatus equipped with a twister-winder. TNFYs with various yarn diameters, twist numbers, twist angles, and bulk densities were produced by varying the twisting speed and the winding speed, and TNFYs with a wide nT range of 395–1276 T m−1 were obtained. The yarn diameter of the TNFY at each winding speed decreased as the twisting speed increased, and the twist angle and bulk density increased with increasing twist speed. The fiber diameter of the individual nanofibers comprising a TNFY depended only upon the laser-irradiation conditions and was independent of twist conditions. The nanofiber diameters of all TNFYs were in the range from 370 to 400 nm and were almost equal because the laser-irradiation conditions were kept constant in the preparation of the TNFYs. In the relationship between the twist angle and the degree of crystal orientation for almost of the annealed TNFYs, the experimental data points (except SW = 0.286 m min−1) fell approximately on a straight line, suggesting that the NFs were firmly twisted by the twister153
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