Core–Shell Fibers Electrospun from Phase-Separated Blend Solutions

Tainan 701, Taiwan, Republic of China. Biomacromolecules , 2017, 18 (9), pp 2906–2917. DOI: 10.1021/acs.biomac.7b00863. Publication Date (Web): ...
3 downloads 15 Views 10MB Size
Download PDF
Article pubs.acs.org/Biomac

Core−Shell Fibers Electrospun from Phase-Separated Blend Solutions: Fiber Formation Mechanism and Unique Energy Dissipation for Synergistic Fiber Toughness Chi Wang* and Ting-Ting Hsiue Department of Chemical Engineering, National Cheng Kung University, Tainan 701, Taiwan, Republic of China S Supporting Information *

ABSTRACT: Through single-tube electrospinning, the biodegradable core−shell fibers of poly(3-hydroxybutyrate) (PHB) and poly(D,L-lactic acid) (PDLLA) were obtained from blend solutions with different compositions at a total polymer concentration of 7 wt %. Regardless whether PHB is the major or minor component (PHB/PDLLA = 90/10, 75/25, 50/50, and 25/75 wt. ratio), these phase-separated solutions all yielded core−shell fibers with PHB as core and PDLLA as shell. A new scenario of core−shell fiber formation was proposed on the basis of the relative magnitude of the intrinsic relaxation rate of fluids and external extension rate during electrospinning. The effects of blend compositions on the morphologies of the Taylor cone, whipping jet, and as-spun fibers were investigated. The diameters of core−shell fibers can be tailored by simply varying the PHB/ PDLLA ratios. Two scaling laws describing the apparent viscosity (ηo) dependence of the outer fiber diameter (dfo) and core fiber diameter (dfc) were derived. That is, dfo ∼ ηo0.38 and dfc ∼ ηo0.86. The microstructures of the as-spun fibers were determined by differential scanning calorimetry, Fourier transform infrared spectroscopy, and synchrotron wide-angle and small-angle X-ray scatterings. Results showed that the PDLLA component was in the amorphous state, and the crystallizability of PHB component remained unchanged, except the amorphous 10/90 fibers electrospun from a miscible solution state. The synergistic mechanical properties of the core−shell fibers were obtained, along with the ductile PDLLA shell enclosing the brittle PHB core. The enhanced toughness was attributed to the fragmentation of the brittle PHB core and necking fracture of the ductile PDLLA shell, which served as an effective route for energy dissipation. Compared with the neat PHB fiber, the 90/10 and 75/25 core−shell fibers possessed larger elastic moduli, which was attributed to the high PHB crystal orientation in their core sections despite the reduced PHB crystallinity. By contrast, the crystal c-axis of PHB in the 25/75 core−shell fibers was preferentially perpendicular to the fiber axis, suggesting the significant stretching of developing PHB crystals during electrospinning.

1. INTRODUCTION

outer tubes to form a composite Taylor cone with an outer meniscus surrounding an inner one.6 At a critical applied voltage, an electrified coaxial jet is issued from the apex of the Taylor cone. After the solvent evaporates during jet bending (whipping), the core−shell fibers are collected on the grounded electrode. The electrospun fibers can be used to encapsulate drugs (such as those containing enzymes, growth factors, and DNA) in the core section and release them in a controlled manner, thereby further advancing the core−shell fibers to

Exhibiting high surface area to volume ratio and high porosity with interconnected networks, biodegradable nanofibrous membranes prepared through electrospinning can be potentially used as scaffolds for tissue engineering.1 For cell culture, electrospun fibers with soft fiber surfaces (or low Young’s moduli) are usually preferred because they can provide an environment conducive for cell growth. Aligned nanofiber scaffolds were successfully used to promote the cell proliferation in the fiber direction.2 Recently, researchers found that core−shell nanofibers can be produced through coaxial-tube electrospinning.3−6 During coaxial-tube electrospinning, two different solutions are delivered to the inner and © XXXX American Chemical Society

Received: June 19, 2017 Revised: August 15, 2017

A

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules many potential applications.1,7−11 Compared with the singlecomponent fibers, core−shell fibers provide a feasible route for a better controlled release of embedded chemicals and drugs7−11 through an adjustment of the core/shell microstructure (e.g., crystallinity, crystal orientation, and fiber diameters). Carefully selecting suitable components for the core and shell enables not only the customization of the drug release kinetics but also the control of the degradation rate, as well as the mechanical properties. Through these changes, variations in scaffold affinity toward the cultured cells can be obtained. Nonwoven scaffolds composed of core−shell fibers with a rigid core and soft shell facilitates the retention of mechanical integrity and provides good contact with the cultured cells. In contrast to the coaxial-tube process, single-tube electrospinning can be used to produce core−shell fibers.12−15 This alternative process is often preferred because of its simple and convenient experimental setup. However, the mechanism that affects the structure of the produced core−shell fibers remains unclear. This mechanism possibly involves the interplay between thermodynamic factors (blend compatibility, polymer solubility, and so on) and kinetic factors (solvent evaporation, phase separation kinetics, and others). In general, the latter is a more dominant factor than the former.12 Poly(3-hydroxybutyrate) (PHB) is a biosynthetic aliphatic polyester used in practice. However, PHB is unsuitable for practical application in biomedical fields because of its brittleness due to its large spherulites. The physical properties of PHB and its biodegradable nature can be improved by reducing its crystallinity in order to increase its flexibility. In this approach, PHB copolymers containing other monomer units, such as hydroxyhexanoate, are produced. Another effective route for properties improvement is simply to blend PHB with another biodegradable polymer, such as poly(lactic acid) (PLA), which is a chemically synthesized aliphatic polyester and possesses better toughness. However, although singlecomponent PHB and PLA nanofibers are extensively addressed in literature,16−18 electrospinning of PHB/PLA blend solutions is not investigated yet. In our previous articles, the relationships between the processing−structure−properties of single-component PHB19 fibers and those of PDLLA20 fibers were explored. In addition, the core−shell fibers of PHB/PDLLA and PDLLA/PHB can be obtained through coaxial-tube electrospinning, along with the interchanging between the PHB and PDLLA solutions in the inner and outer tubes, respectively.21 As-prepared core−shell fibers demonstrate two-stage release kinetics when drugs are loaded in the core section. A burst release of a small amount of drug occurs in the first stage, followed by the zero-release order of the remaining drugs in the second stage. As a series study on biodegradable fibers, the present work aimed at preparing core−shell fibers through single-tube electrospinning of PHB/ PDLLA blend solutions. Chloroform (CF) is a good solvent for both PHB and PDLLA. However, the one-phase CF blend solution of PHB/PDLLA at high temperatures (∼70 °C) became phase separated at ambient conditions. Moreover, the electrospinning of the CF solutions is generally difficult to perform because CF solution has low solution conductivity and high solvent volatility. In this work, to improve the spinnability of the solution, we used a solvent mixture of CF and dimethylformamide (DMF), which did not dissolve PHB, but dissolved PDLLA at 25 °C. Through this approach, we were able to elicit nonsolvent-induced phase separation in the PHB-

rich phase. The electrospinning of the phase-separated blend solutions with different compositions was then performed, and the as-spun nonwoven fabrics were characterized. Results indicated that the core−shell fibers with PHB cores were obtained even from systems with PHB-rich solutions (PHB/ PDLLA 90/10 and 75/25). Therefore, core−shell fibers with hard PHB core and soft PDLLA shell may be modified for tissue engineering. More importantly, the toughness of PHB/ PDLLA 50/50 nonwoven fabrics is nearly 1 order of magnitude higher than that composed of neat PHB fibers. On the basis of fracture surface, a unique energy dissipation mechanism was proposed to account for this synergistic mechanical strength.

2. EXPERIMENTAL SECTION Solution Preparation and Properties. PDLLA pellets and PHB powders were purchased from Natureworks LLC and Sigma-Aldrich, respectively. For the PDLLA, the molar content of D-lactide units was 10%, and the weight-average molecular weight (Mw) and polydispersity index were 1.78 × 105 g/mol and 2.1, respectively, determined from gel permeation chromatography. Using a capillary viscometer, the viscosity average molecular weight of PHB was determined to be 2.62 × 105 g/mol. Chloroform is a good solvent for both PDLLA and PHB. However, its conductivity (dielectric constant) and boiling temperature are low, which may hamper processing during electrospinning. Thus, a solvent mixture of CF and DMF with a weight ratio of 9/1 was used to prepare PHB/PDLLA blend solutions in order to enhance solution conductivity and reduce solvent volatility. DMF is a good solvent for PDLLA but is a nonsolvent for PHB at room temperature. The PHB/PDLLA blend solutions with a fixed polymer content of 7 wt % but different weight ratios (100/0, 90/10, 75/25, 50/50, 25/75, 10/90, and 0/100) were prepared for electrospinning. Solution properties, including surface tension, conductivity, and shear viscosity, were measured using the Face surface tension meter (CBVP-A3), Consort conductivity meter (C832), and Brookfield viscometer (LVDV-I+, spindle 18, and cup 13R), respectively. Electrospinning Process. A needle with an inner and outer diameter of 1.07 and 1.47 mm was used as the spinneret. The prepared solution was delivered by syringe pumps (Cole-Parmer) to the needle at a controlled flow rate. A high potential was applied to the spinneret by a high-voltage source (Bertan, 205B) to provide a sufficient electric field for electrospinning. To construct a needle-to-plate electrode configuration, a steel net (30 × 30 cm2) was used as collector for the electrospun fibers at a fixed tip-to-collector distance of 14 cm. A gas jacket was used to introduce CF vapors using N2 as a carrier gas to appropriately encapsulate the Taylor cone without interrupting jet bending during electrospinning. The morphologies of the electrified straight jet were monitored to derive the following: (i) the location for jet whipping, Lj, which is measured from the needle end to the initiation of jet whipping, and (ii) the terminal jet diameter at the end of the straight jet, dj, which was determined by a laser diffraction technique. The experimental details were described in a previous article.19,20 For the present needle-to-plate configuration setup, the electric field calculation was performed using the FLUX2D9.10 software to solve the 3-D electrostatics problem. The electric field intensity at the initial whipping position (z = Lj) was then determined and denoted to be Ej.20 In addition to nonwoven type fibers, aligned fibers were collected using a rotating drum with its surface covered by thin metal wires as a grounded collector.22 The linear velocity of the rotating drum was 3.8 m/s. Aligned fibers were used to characterize the global chain (amorphous plus crystalline) orientation within fibers by polarized infrared (polarized IR), crystal orientation by wide-angle X-ray diffraction (WAXD), and lamellar morphology by small-angle X-ray scattering (SAXS). Fiber Characterization. The morphology of the fibers was observed under a scanning electron microscope (SEM, Hitachi S4100). The fiber diameters were measured from a collection of ∼500 fibers, from which the average fiber diameter was determined. B

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules

Figure 1. Phase-contrast micrographs of the 7 wt % PHB/PDLLA 50/50 blend solution after being cooled from 70 to 25 °C. (A, left column) in CF solvent and (B, right column) in CF/DMF 9/1 cosolvent. Phase separation of the blend solution is observed in both solution systems. In (A), it is attributed to TIPS to develop the cocontinuous structure of PHB-rich and PDLLA-rich phases. In (B), it is attributed to TIPS as well as NIPS. The DMF-induced phase-separated structure is disclosed as tiny irregular particles with a diameter of ∼1 μm dispersed in the PHB-rich domain. The thermograms of the as-spun fibers were obtained by using a differential scanning calorimeter (PerkinElmer, DSC7) under a nitrogen atmosphere at a scanning rate of 10 °C/min. For the melting endotherm measured, the peak temperature (Tm) and enthalpy (ΔHm) were recorded. The fiber crystallinity (ϕDSC c ) was determined using the ratio of ΔHm/ΔH°m, where ΔH°m is the melting enthalpy for a 100% crystal. The corresponding ΔH°m value is 146 J/g for the PHB with pure α-form crystallites.23 The porosity of fiber membranes was calculated from the density measurement according to the procedure suggested by Zong et al.24 The mechanical properties of the fiber membranes were obtained from the stress−strain curve using a universal tensile testing machine (Instron 4500) at a stretching rate of 10 mm/min. Specimens with dog-bone shape were prepared using a sharp cuter. The mean thickness of each sample was ∼35 μm. The gauge length and sample width were 15 and 4 mm, respectively. The reported data of mechanical properties represent the average results of the three tests. The Fourier transform infrared (FTIR) spectra of the composite fibers were obtained in a PerkinElmer Spectrum 100 spectrometer with a resolution of 2 cm−1 and capacity of 32 scans. A typical FTIR spectrum in the range of 1650−1850 cm−1 was used for analysis, in which the characteristic band of the carbonyl group was 1757 cm−1 for the amorphous PDLLA.25 Meanwhile, the bands at 1722, 1728, and 1742 cm−1 were attributed to the crystalline, interfacial, and amorphous phases of the PHB component, respectively.26 Following the previous model analysis,26 the deconvolution of the obtained spectrum was successfully performed to separate the contribution of both PDLLA and PHB components. To monitor the changes in chain orientation as a consequence of processing variations, the aligned fibers were evaluated to calculate the Hermans orientation function (f IR). Dichroic ratio (R) was calculated using the formula R = A∥/A⊥, where A∥ and A⊥ are the absorbances with the polarized IR beam, parallel and perpendicular to the fiber direction, respectively. After deconvolution of the A∥ and A⊥ spectra, the 1756 cm−1 band yielded the R value for the PDLLA component. Furthermore, the other three bands produced the R value for the PHB component. The Hermans orientation function can be correlated with the measured R by27

fIR = [(R o + 2)/(R o − 1)][(R − 1)/(R + 2)]

where Ro = 2 cot2 ψ, with ψ being the angle between the transition moment vector and the chain axis. For the carbonyl groups, ψ can be assumed to be 90°. Depending upon the main chain orientation, the value of f IR is between −0.5 and 1.0. The lower limit value represents the chain alignment in the plane perpendicular to the fiber direction, whereas the upper limit is for the chains orienting perfectly parallel to the fiber axis. The value f IR = 0 corresponds to the random chain orientation in the fiber. 2-D SAXS and 2-D WAXD experiments on the aligned fibers were conducted at the beamlines BL23A and BL17A at the National Synchrotron Radiation Research Center in Taiwan, respectively. The contribution from air scattering was subtracted from the 2D WAXD and SAXS patterns. The details of the experimental setup and data analysis can be found elsewhere.28 For the 2-D SAXS experiment, the aligned fibers were sealed in Al pans with Kapton windows ca. 2 mm in diameter for the X-ray beam. With an 8 keV (wavelength = 1.55 Å) beam, the SAXS data were collected by 2-D (200 × 200 mm2) proportional counters in a master-slave mode. The scattering vector q (=4π sin θ/λ, where 2θ is the scattering angle and λ is the wavelength of X-ray) ranged from 0.01 to 0.2 Å−1 for SAXS. The SAXS intensity profile along the fiber direction was used to deduce the long period (L) and lamellar thickness (lc) through the 1-D correlation function analysis.29 Prior to the correlation analysis, the raw intensity (Iraw) at high q regions was extrapolated with the aid of Porod law:

lim [K p − (Iraw − Ifl)q 4] = 0

q →∞

(2)

where Ifl is the thermal density fluctuation, and Kp is the Porod constant. Because of the strong zero-angle scattering,30 the intensity at low q regions near the beamstopper also had to be modified to extract contribution exclusive from the lamellar morphology. Debye−Bueche model that describes intensity scattered by the random distribution of heterogeneity was applied to fit the results at low q regions. The SAXS intensities, corrected by the Porod law and Debye−Bueche model, took the Fourier transformation to derive the corresponding correlation function, from which L was obtained from the first maximum, and lc was determined according to the work of Strobl et al.29 For the WAXD experiment, a MAR345 2-D detector of 380 × 250 mm2 imaging area was used. The azimuthally integrated profiles of

(1) C

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules WAXD were extracted from the 2-D patterns, from which the squared average of cosine value (⟨cos2 ϕhkl⟩) of the angle, ϕhkl, between the normal of (hkl) reflection with the fiber axis, was evaluated by

cos2 ϕhkl

∫0

=

∫0

π /2

mixture of CF/DMF 9/1 (w/w) was used in this study to enhance solution conductivity and reduce solvent volatility. Through this approach, a stable cone-jet electrospinning mode was ascertained for hours, thereby yielding reproducible fibers for comparison. However, the prepared solutions may be phase-separated at ambient temperature because of TIPS and NIPS. Figure S1 shows the appearance of prepared solutions prior to electrospinning. After the hot solutions were cooled down to ∼25 °C (time = 0), the blend solutions of PHB/ PDLLA 90/10, 75/25, 50/50, and 25/75 were turbid and phase-separated into PHB-rich domain and PDLLA-rich domain. The former was opaque, whereas the latter became transparent. Meanwhile, the neat PDLLA solution and 10/90 blend solution were in a single phase and were transparent. After 120 h, the 10/90 blend solution was also phase-separated, forming the cotton-like PHB-rich phase dispersed in the PDLLA-rich phase matrix. After 30 days, macroscopic gels were seen for the solutions of 100/0, 90/10, 75/25, and 50/50 by inverted vial method. This finding suggested that the PHB-rich solutions were able to form the gel structure. Our observation is consistent with a previous report that the gelation of PHB in DMF proceeds after the thermally induced liquid−liquid phase separation.33 By contrast, the neat PDLLA solution remained in a single phase after 30 days. The detailed discussion for the phase separation kinetics is beyond the scope of the present study and will be reported in our forthcoming article. For fair comparison, all the measurements of the solution properties and the performance of electrospinning were carried out within 5 h after the prepared solutions were cooled down to ambient temperature. The composition dependence of solution viscosity (ηo) is shown in Figure 2. The neat PHB solution had

Ihkl(ϕ)cos2 ϕ sin ϕdϕ/

π /2

Ihkl(ϕ)sin ϕdϕ

(3)

where Ihkl(ϕ) is the intensity distribution of (hkl) reflection on the Debye ring, and ϕ is the azimuthal angle measured from the meridian (fiber axis direction). The orientation of induced crystals was determined using the Herman orientation function as well:31

fhkl = (3 cos2 ϕhkl − 1)/2

(4)

Both the f hkl deduced from WAXD and the f IR (global chain orientation, including the amorphous and crystalline phases) estimated from the polarized IR were crucial in determining the mechanical properties of the electrospun fibers.

3. RESULTS AND DISCUSSION For a given polymer solution, the determination of entanglement concentration (ϕe) is an important task prior to electrospinning. Solutions with concentrations lower than ϕe generally yield electrospun products with beads-on-a-string structures because of capillary instability. Based on the log−log plot of the solution specific viscosity versus its volume fraction, the ϕe for the PHB/(CF/DMF) solution was determined to be 5 wt %.19 For bead-free fiber production, a minimum concentration of 7 wt % is required for PHB/(CF/DMF) solution.19 Thus, we fixed the polymer content to 7 wt % but varied the PHB/PDLLA compositions to prepare the blend solutions for electrospinning. Moreover, our previous study also showed that the ϕe for the PDLLA/DMF solution is 10 wt %, and the minimum concentration to produce bead-free fibers is 20 wt %.20 Thus, the 7 wt % PDLLA/(CF/DMF) solution was in the semidilute solution state in the absence of chain entanglements. Beaded PDLLA fibers were produced when the solution was electrospun. Solution Properties and Solution Phase Behavior. The high molecular weight blends of PHB and PLA are immiscible in the melt state.32 Despite that CF is a good solvent for both PHB and PLA, the CF solutions of PHB/PDLLA blend exhibited a single phase at high temperatures but became phase-separated when cooled at 25 °C. Figure 1A demonstrates the phase evolution of the CF solution with PHB/PDLLA 50/ 50 blend subjected to a temperature jump from 70 to 25 °C. Spinodal phase separation took place at 25 °C, and the coarsening of the interconnected domains was observed due to the thermally induced phase separation (TIPS). Meanwhile, nonsolvent-induced phase separation (NIPS) was also perceived in the originally single-phase 7 wt % PHB/CF solution at 25 °C by introducing DMF nonsolvent when the CF/DMF wt. ratio was 92/8. Although DMF is completely miscible to CF, the addition of DMF caused NIPS because DMF is insoluble to PHB. For the CF/DMF 9/1 solution of 50/50 polymer blend, both TIPS and NIPS played a role in the liquid−liquid phase separation, as shown in Figure 1B. In the PHB-rich domains, many tiny irregular particles induced by the DMF nonsolvent were clearly seen. Provided that neat CF solvent is used to prepare the PHB/ PDLLA blend solutions for electrospinning, reaching a stable process for continuous fiber production is difficult because of the low conductivity and high volatility of CF. Thus, a solvent

Figure 2. Viscosity of PHB/PDLLA solutions with different compositions. The solid content of polymer in the CF/DMF cosolvent is 7 wt %. The inset shows the conductivity and surface tension of the corresponding solutions.

a ηo of 580 cP, which was about 4× larger than that of the neat PDLLA solution. An increase in the PDLLA content led to an apparent decrease in solution viscosity. The ηo of PHB/PDLLA 10/90 solution was similar to that of the neat PDLLA solution. The composition dependence of solution conductivity (κ) and surface tension (γ) was also provided in the inset. The κ of cosolvent was 1.6 μS/cm. The measured κ values for the neat PDLLA and PHB solutions were 1.6 and 2.1 μS/cm, respectively. A shallow minimum κ of 1.4 μS/cm was observed for the PHB/PDLLA 25/75 solution. Meanwhile, the magnitude of γ remained relatively constant at 27.1 dyn/cm, regardless of blend compositions. Based on these measureD

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules ments, the main difference in the electrospinning solutions was solution viscosity rather than conductivity or surface tension. Core−Shell Fiber Discovery. Electrospinning with a stable cone-jet mode is favorable for obtaining reproducible results for fair comparison. At a given flow rate (Q) and tip-to-collector distance, the range of applied voltage (V) for the cone-jet mode is rather narrow and different from one solution to another. Therefore, the construction of processing window (V−Q plot) is essential to determining the common processing variables for the prepared blend solutions, as shown in Figure 3. At a fixed

Figure 3. Processing window for the electrospinning of the 7 wt % solutions with different PHB/PDLLA compositions. Filled and open symbols stand for the lower and upper bound voltage, respectively, required to develop a stable cone-jet electrospinning mode at a given Q. The tip-to-collector distance is fixed at 14 cm. The asterisk represents the common Q and V applied to electrospin the blend solutions.

Q, the filled and open symbols are the lowest and highest voltages, respectively, required to obtain a stable cone-jet mode for electrospinning. Only a small range of applied voltage was available for operation. At a given Q = 3 mL/h, we found that 7.4 kV was applicable for all the blend solutions at a fixed tip-tocollector distance of 14 cm. By applying these processing variables, a stable cone-jet spinning mode was ascertained, and the dimensions of cone height (Hc), jet length (Lj), and straight jet-end diameter (dj) were measured. The morphology of the as-spun fibers is shown in Figure 4A, where uniform and beadfree fibers were produced from all the blend solutions. The measured fiber diameters (dfo) of the as-spun neat PHB, 90/10, 75/25, 50/50, 25/75, and 10/90 composite fibers, were 2.5 ± 0.4, 2.9 ± 0.3, 2.7 ± 0.4, 2.5 ± 0.2, and 1.9 ± 0.3 μm, respectively. Increasing the PDLLA content in the blend solution led to the reduction in the diameters of the composite fibers. The as-spun PHB/PDLLA fibers may have different fiber morphologies, such as matrix-dispersed, cocontinuous, or core−shell structures.34 The as-spun fibers were immersed in the DMF solvent with vigorous shaking for an extensive period to extract all the PDLLA components and reveal the fiber morphologies. The complete removal of PDLLA component was confirmed by a comparison between FTIR spectrum before the solvent treatment and that after (Figure 5c). The extracted fibers were then observed again. The corresponding SEM images are shown in Figure 4B. Continuous fibers with remaining PHB component were observed in the 90/10, 75/25, 50/50, and 25/75 electrospun fibers. This finding suggested the formation of core−shell structures in the as-spun fibers containing PHB in their core sections and PDLLA in their shell sections. The diameter of core fiber (dfc) was measured. The dfo and dfc values decreased with increasing PDLLA

Figure 4. SEM images of as-spun fibers obtained from solutions with different compositions of PHB/PDLLA (A, left column) and fibers after DMF extraction of the PDLLA component (B, right column).

composition in the spun solution. By contrast, the preservation of residual PHB fibers for the 10/90 fiber was not observed after the removal of the PDLLA component. Instead, a PHB thin film with gel-like structure was observed after long immersion in DMF. Therefore, the as-spun 10/90 fibers did not form the core−shell structure. To further verify the core fiber diameter dfc of the core−shell fibers, an indirect method was also used. Provided that the outer fiber diameter dfo was measured, the value of dfc can be estimated by the following equation: dfc = dfo[ϕv′/(1 + ϕv′)]0.5

(5)

where ϕv′ is the volume ratio of the core−shell component (Vc/ Vs) in the as-spun fiber, which can be determined by FTIR. The typical FTIR spectrum of an as-spun fiber membrane is shown by curve a in Figure 5a, where a pronounced curvature of the baseline at the high wavenumber region is observed, mainly due to the noneven thickness of samples. Notably, a downwardpointing feature near the desirable wavenumber region of the E

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules

cm−1 were relevant to the crystalline, interfacial, and amorphous phases of the PHB, respectively. The amorphous state of the PDLLA component was confirmed from the WAXD intensity profile. Following the previous model analysis,26 the deconvolution of the obtained spectrum was successfully performed, and thus, the contribution of the PHB and PDLLA components (bottom curves) were distinguished. Under the assumption that the absorption coefficient of these CO stretching vibration bands are similar,35 the ratio of the carbonyl group content, X′, of the as-spun core−shell fibers was derived. The volume ratio of ϕv′ can be further derived after obtaining the value X′, because each PHB and PDLLA monomer contains a single carbonyl group. The calculated ϕ′v value provided the measured composition of the as-spun fibers. Thus, eq 5 can be revised to dfc = dfo[X ′Mo′ /(ρ′ + X ′Mo′)]0.5

(6)

where M′o and ρ′ are the monomer weight ratio and the solid density ratio of the core to shell components, respectively. Figure 6 shows the ηo-dependence of jet diameter dj, outer fiber diameter dfoc, and core fiber diameter dfc. In addition to dfc

Figure 6. Viscosity dependence of the jet diameter dj measured by light scattering (□), outer fiber diameter dfo measured by SEM (△), and core fiber diameter dfc measured directly by SEM (○) and derived indirectly from FTIR results (▽). The filled symbols are for the neat PHB. The weight ratios of PHB/PDLLA for solutions with ηo from low to high as follows: 10/90, 25/75, 50/50, 75/25, and 90/10. The shaded region is for the core/shell fibers.

Figure 5. (a) FTIR spectra of an as-spun 50/50 fiber mats (curve a), and fiber mats filled with PDMS fluid (curve b) and mineral oil (curve c). The arrow shows the presence of an absorption dent. (b) Deconvolution of the FTIR spectrum of 50/50 fibers to determine the relative contribution of each band. The solid line and dashed line of the upper curves are the measured data and convolution of the individual bands (bottom curves). The band at 1756 cm−1 is related to the PDLLA, whereas the bands at 1742, 1728, and 1721 cm−1 are associated with the PHB chains in the amorphous, interfacial, and crystalline region, respectively. (c) FTIR spectra before and after removal of the PDLLA component. A complete removal of PDLLA component is observed through the disappearance of the 1756 cm−1 band.

measured directly under SEM after PDLLA extraction (Figure 4B), the values of dfc estimated by eq 6 are also included. Fairly consistent dfc values are observed thereby further verifying the formation of core−shell fibers. More importantly, eq 6 provides a convenient estimate of the core fiber diameter without using the complicated fluorescence dye technique, which is frequently applied to the core−shell fibers with a fiber diameter of several micrometers. For the limited ηo range available (165−500 cP), the data between ηo and dj were relatively scattered. The blend solutions producing core−shell fibers had a similar dj value of 11.9 ± 0.4 μm, which was lower than that for the PHB/PDLLA 10/90 solution (14.8 ± 0.4 μm). Moreover, both the dfo and dfc exhibited good correlations with ηo, and the following scaling laws were derived: dfo ∼ η0.38 with a R2 coefficient of 0.979 and o dfc ∼ η0.86 with a R2 value of 0.963. Despite the usage of the o blend solution of PHB/PDLLA, the present result showed that the derived exponent for the dfo−ηo relation is similar to the pure PHB fiber (∼0.41)19 and pure PDLLA fibers (∼0.45).20 A previous study showed that a short straight jet with a small dj produces solid fibers with small diameters.19 Therefore, dfo is relevant to straight jet-end diameter dj and the jet whipping process that subsequently occurs. The effectiveness of jet

carbonyl group (1650−1850 cm−1) was observed. This feature was attributed to the coupling of the real and imaginary parts of the refractive index caused by scattering, resulting in a slightly dispersive line shape. Several refractive-index matching liquids were used to reduce abnormal scattering in order to improve the quantitative analysis results. Apart from polydimethylsiloxane (PDMS) used previously,21 mineral oil can be also applied to resolve this problem. As shown in Figure 5a, the FTIR spectrum of the mineral oil-wetted membrane (curve c) exhibited an even baseline, and the absorbance dent at 1790 cm−1 disappeared. The PHB/PDLLA amount in the core−shell fiber can be achieved by analyzing the FTIR spectra. Figure 5b shows a typical FTIR spectrum (range: 1650−1850 cm−1) of the as-spun fibers obtained from the 50/50 solution. The C O stretching vibration band at 1757 cm−1 was associated with the amorphous PDLLA, and the bands at 1722, 1728, and 1742 F

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules

Table 1. Effects of Polymer Compositions on the Jet Length (Lj), Electric Field at Straight Jet End (Ej) during Processing, As Well As the Physical Properties of As-Spun Fiber Mats; The Glass Transition of PDLLA Component (Tg,PDLLA), the Melting Temperature of PHB Component (Tm,PHB), the PHB Crystallinity Measured by DSC and FTIR (ϕDSC and ϕFTIR ), and the c c a Porosity of Nonwoven Mats PHB/PDLLA 100/0 90/10 75/25 50/50 25/75 10/90

Lj (mm)

Ejb (MV/m)

± ± ± ± ± ±

6.5 6.0 7.3 6.7 6.6 33.6

26.0 27.0 24.5 25.6 25.7 11.6

0.8 0.8 0.7 0.9 0.7 0.5

Tg,PDLLA (°C)

Tm,PHB (°C)

ϕDSC c

ϕFTIR c

porosity (%)

51.8 51.3 51.9 46.6 45.2

169.9 170.4 169.5 167.4 166.9 166.4

0.61 0.52 0.42 0.29 0.14 0.05c

0.62 0.53 0.45 0.34 0.17 ∼0

63 65 66 66 66 68

a The italic area is for the core−shell fibers. bEj is the electric field at the straight jet end calculated from numerical FEA analysis using the following dimensions; the inner and outer diameters of the needle spinneret are 1.07, and 1.47 mm, respectively, and the needle length is 40 mm. cNot the crystallinity of the as-spun fibers, but the crystallinity of the thermally induced PHB crystals during DSC heating scan.

deform under a high extension rate (∼102−103 s−1) induced by the setup electric field remains challenging. Domain deformation in the two-phase solution systems subjected to a shear field is extensively studied.39 On this basis, the phase-separated domain morphology in the electrospinning jet, which is mainly subjected to an extensional field, depends on the relative magnitude of the external extension rate (ε̇) and the relaxation rate of interfacial deformation (τ−1 s , where τs is the characteristic relaxation time of the dispersed droplet expressed by ηoR0/σ, with R0 as the droplet radius and σ as the interfacial tension). τs−1 is rather small because of weak interfacial tension. Droplets are dispersed into small droplets that are elongated parallel to the flow field at intermediate extension rates (ε̇ > τ−1 s ). This kind of domain deformation possibly occurs in the spinneret tube and in the Taylor cone after emerging from the spinneret tube. At higher extension rates encountering in the Taylor-cone apex, the elongated domains with a large aspect ratio rapidly transform to become strings. When the extension rate in the spin line is too high (i.e., ε̇ > 102 s−1), the polymer chain relaxation within the phaseseparated domains may be eventually affected. At this stage, the relaxation rate of polymer chains in the PHB-rich phase (τ−1 PHB) and that in the PDLLA-rich phase (τ−1 PDLLA) must be considered. Previous studies showed that the entanglement concentrations for the related PHB and PDLLA solutions are 5 and 10 wt %, respectively.19,20 In the present study, for the 7 wt % electrospinning solution, the PHB-rich phase was observed in the entangled regime, but no chain entanglement was detected in the PDLLA-rich phase. Moreover, gelation occurred in the PHB-rich phase because of NIPS, which further elevated the τPHB value. These considerations intuitively led to the following −1 −1 relation: τ−1 PDLLA > τPHB ≫ τs . On this basis, a more effective chain extension occurs in the PHB-rich domain than in the PDLLA-rich domain, regardless of whether PHB is the majority or minority phase. Under a strong extensional flow (say, τ−1 PDLLA > ε̇ > τ−1 PHB), within the elongated PHB-rich phase the string-like structures are likely to develop and squeeze out the solvent through the interface to further dilute the adjacent PDLLA-rich phase. The developing string-like structures tend to migrate toward the center of the tapering jet because of the circumferential stress induced by the uniaxial flow field. The migration and coalescence of the PHB-rich strings expel the less viscous (and less deformed) PDLLA-rich phase to the jet surface, thereby producing the core−shell jet, which eventually produces the core−shell fiber after solvent evaporation. Multicore−shell fibers can be produced provided that the viscosity of PDLLA-rich phase remains high to obstruct the

whipping is relevant to the electric field at the straight jet end. As the electric field decays exponentially with the distance from the needle end,19 a straight jet with a small Lj is subjected to a high electric field Ej that, in turn, produces thin fibers. As shown in Table 1 and Figure 6, both the Lj and dj remained relatively unchanged until the PDLLA content reached 90%. The PHB/ PDLLA 10/90 solution produced a short jet with a large dj. The significant reduction of Lj led to an enhanced Ej (ca. 33.6 MV/ m vs 6.5 MV/m for the others) for the jet whipping to effectively reduce the fiber diameter, despite the large dj. Therefore, the jet geometry Lj plays a more important role than dj in determining the final fiber diameter. Although several attempts were performed to predict the Lj value during electrospinning,36 present efforts remain unsatisfactory because of the complex interaction between the solution properties and processing variables. More importantly, the genuine mechanism that triggers (or initiates) the jet whipping process is unknown. Proposed Mechanism for Core−Shell Fiber Formation. Regardless of PHB/PDLLA compositions (90/10, 75/25, 50/50, and 25/75 wt. ratio), the phase-separated solutions all yielded core−shell fibers with PHB cores. The formation of core−shell fibers by single-tube electrospinning of phaseseparated solution is attracting considerable attention since the pioneering studies of some research groups.12−14 However, the detailed kinetics pathways leading to the core−shell formation are not fully elucidated. An early argument was based on consideration of the system free energy.12,37 This argument was learned from the steady-state tube flow system for twocomponent blend melts with a confined wall to produce spatial regions with different shear rates. It inferred that the phaseseparated phase with high viscosity migrates toward the tube center region with low shear rate to reduce the total free energy of the system.37 At first glance, this argument seems to apply to the present solution system because the viscosity of PHB-rich phase is higher than that of the PDLLA-rich phase (Figure 2). However, this hypothesis should be carefully examined because it applies the principle of minimizing the free energy in nonequilibrium states. Electrospinning is a fast and transient process that produces fibers within several milliseconds, and the kinetic (or dynamic) factor is generally assumed to dominate the thermodynamic factor (such as surface tension38). Moreover, the extension rate of the electrospinning jet may dramatically vary along the spin-line (from the Taylor cone to straight jet and whipping jet) but is more or less uniform in the cross-section of the thin jet because of the free-surface flow field. Therefore, determining how the phase-separated domains G

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules string-like domain migration and/or coalescence.14 Our argument is based on the previous study,40 in which the extension rate of the electrospinning jet was found to be higher than the relaxation rate of the polymer chains in the singlephase polymer solution, plausibly leading to the flow-induced phase separation in the strong extensional flow. Microstructure of Core−Shell PHB/PDLLA Fibers. The WAXD patterns of electrospun fibers are shown in Figure 7.

Figure 8. DSC heating traces of the as-spun fibers with different PHB/ PDLLA compositions. The dashed line shows the baseline to distinguish the cold crystallization event (75−115 °C) from the melting behavior of PDLLA crystals (with a melting peak at about 124 °C) developed during scan-annealing process.

variations among the thermal properties of the core−shell fibers. The measured data are tabulated in Table 1. A significant Tg drop in PDLLA component was observed from 57 °C for the neat PDLLA fiber to 45 °C for the 10/90 fibers. Further increase in the PHB component gradually increased the Tg of PDLLA component to a value of 51−52 °C in the core−shell fibers. For the core−shell fibers, no cold crystallization occurred. The determined PHB crystallinity ϕDSC decreased c at increased PDLLA content. This result is in agreement with those determined from deconvolution of the FTIR spectra, that is, ϕcFTIR = A1722cm−1/(A1722cm−1 + A1728cm−1 + A1742cm−1). Moreover, the crystallizability of PHB component (the value of ϕDSC normalized by the weight fraction of PHB component) c remained relatively unchanged for the core−shell fibers derived from the 90/10, 75/25, 50/50, and 25/75 blend solutions. A cold crystallization peak was observed for the 10/90 composite fibers at 94.9 °C, which was lower than that for the neat PDLLA fiber. In addition, the measured crystallization enthalpy of 5.5 J/g was larger than that for the neat PDLLA fibers (∼1.4 J/g). These results indicated that PHB crystallization is also involved during heating and facilitates the PDLLA crystallization at low temperatures because of the reduction in Tg. The thermally induced crystalline structure in the annealed 10/90 fibers generated two melting peaks at 124.8 and 166.4 °C. The former is related to the PDLLA crystals and the latter is for the PHB crystals. Mechanical Properties of Core−Shell Fibers. The stress−strain curves of the fiber mat electrospun from the solutions with different compositions are shown in Figure 9. The elongation at break for the neat PHB fibers was (∼8%)

Figure 7. WAXD intensity profiles of as-spun fibers with different compositions of PHB/PDLLA. The reflection peak at 2θ ∼ 20° is associated with the β crystalline form of PHB, and the reflection peaks relevant to the helical α crystalline form of PHB are indexed.

The as-spun PDLLA fiber is amorphous. For the neat PHB fibers, the peaks at 2θ values of 13.4, 16.9, and 21.4−22.5°, which can be indexed as (020), (110), and (101)/(111) reflections, agree well with the values reported by Yokouchi et al.41 for the α-form of the PHB crystal having an orthorhombic unit cell with dimensions (a = 0.576 nm, b = 1.320 nm, and c = 0.596 nm). This unit cell structure consists of two left-handed 21 helices. In addition, a broad reflection at 2θ of 20° is discernible, which is relevant to the β-form crystal with planar zigzag conformation. The formation of β-form was suggested to emanate upon drawing from the amorphous domains between 21 helix α-form lamellae.42 For our electrospun fibers, such formation implies that during solvent evaporation the α-form lamellae are likely to develop first, and a further stretching of nearly dried fibers significantly facilitate the formation of βform crystal. As for the composite PHB/PDLLA fibers, the reflection peaks were reduced with increasing PDLLA content. No reflection peak was detected for the as-spun 10/90 fiber, indicating that it is in the amorphous state. For the SAXS patterns (data not shown), a small intensity hump was observed for the neat PHB fiber and PHB/PDLLA fibers with PHB content of up to 50 wt %. For fibers with the 25/75 composition, no SAXS peak was detected. The SAXS peak was associated with the lamellar packing of the PHB component in the core section because the PDLLA component was amorphous. Based on Bragg’s law, the long period was estimated from the peak position (qm) of the Lorentz-corrected intensity profiles (Iq2−q plot): LB = 2π/qm. The calculated LB value is about 5.3 nm, regardless of the fiber composition. Figure 8 shows the DSC heating curves of as-spun fibers. For the beaded fibers of PDLLA, a pronounced enthalpy recovery was observed at 62.6 °C, followed by a shallow exothermic curve with a peak temperature at 97.0 °C, indicating that cold crystallization occurred. The melting temperature of the thermally induced PDLLA crystals was 124.2 °C. The presence of enthalpy recovery peak suggested that an extra energy was required to relax the tight packing of amorphous PDLLA chains. The addition of PHB component resulted in the

Figure 9. Stress−strain curves of PHB/PDLLA nonwoven mats. The nonwoven mat composed of PHB/PDLLA 10/90 is uniform composite fibers and not core−shell fibers. H

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules lower than that for the neat PDLLA fibers (∼57%), indicating the brittleness of the PHB nature. The observation of the high elongation at break for most of the PHB/PDLLA composite fibers is important; the elongation at break for the 50/50 fibers is about 2× larger than that for the neat PDLLA fibers. For a quantitative comparison, Young’s modulus was determined from the initial slope at small strains, and toughness (strain energy density) was calculated from the integration area of stress−strain curve; the results are plotted against the PHB content in the fibers, as shown in Figure 10. The neat PHB

Figure 11. (a, b) SEM images of PHB/PDLLA 50/50 fibers after tensile test to reveal the ductile shell and brittle core characters, (c) a mechanism to describe the resultant fractured fibers. The fracture initially occurs in the PHB core due to its brittleness. Further stretching results in the neck formation at the fragmenting point of PHB core due to the Poissonian contraction of the ductile PDLLA shell. The neck propagates along the fiber axis, accompanying with a gradual reduction of neck diameter.

shell fibers is established and attributable to the incompatibility between PHB and PDLLA. Further subjecting the fibers to stress leads to the formation of PDLLA necking at the fragmenting points because of the Poissonian contraction of the PDLLA shell. Given that PDLLA is amorphous and ductile, the neck readily propagates along the fiber direction between two PHB fragments until its final failure. This energy dissipation mechanism accounts for the enhanced toughness. Moreover, each long electrospun core−shell fiber provides many sites for neck formation, which further benefits the energy dissipation process. A schematic for this fracture mechanism is provided in Figure 11c. Without sacrificing its Young’s modulus (stiffness), the present core−shell fibers with a rigid core encased by a ductile shell exhibit excellent toughness because of the energy dissipation facilitated by necking fracture. In general, the following major factors must be considered in determining the Young’s modulus of nonwoven mats: (1) fiber diameter, (2) fiber orientation, (3) porosity of the mats, (4) fiber crystallinity, and (5) crystal orientation in the fiber. Among these factors, the obtained PHB/PDLLA fibers have similar fiber diameter, fiber orientation, and porosity (Table 1). Given that 90/10 and 75/25 fibers have lower fiber crystallinities than neat PHB fibers, they are expected to possess a low Young’s modulus when crystal orientations in these fibers are uniform. By contrast, the measured moduli of 90/10 and 75/25 fibers were higher than those of the neat PHB fibers, as shown in Figure 10a. The crystal orientations of PHB in the as-spun fibers merit further investigation. This is discussed in the following section. Crystal Orientation in Core−Shell Fibers. Aligned fibers were collected using a rotating wire drum provided that the linear velocity of the drum is higher than the flight speed of the electrospun fiber to the collector. Figure S2 shows the SEM images of aligned fibers with different compositions. All the fibers exhibited similar fiber orientation in the collecting direction. The outer fiber diameters were measured and are

Figure 10. (a) Young’s modulus and (b) toughness of PHB/PDLLA fiber mats. The shaded area is the blend compositions to produce core/shell fibers.

fibers were stiffer but less ductile than the neat PDLLA fibers. For the structured core−shell fibers, the synergistic effects of the Young’s modulus and toughness were observed. In addition, the two-component fibers may exhibit mechanical properties superior to those of single-component fibers. Compared with neat PHB fibers, 90/10 core−shell fibers had higher Young’s modulus value (166 ± 11 vs 206 ± 10 MPa, respectively). Meanwhile, 50/50 core−shell fibers is tougher than neat PDLLA fibers (7.9 ± 1.1 vs 2.1 ± 0.3 MJ/m3, respectively). After the tensile tests, the fractured fibers were examined under SEM. Typical SEM images are shown in Figure 11a,b for the 50/50 core−shell fibers. Two morphological features merit more discussion. The first feature is the sausage-like texture along the unbroken fiber segment, and the second is the elongated shape at the breaking point. The core−shell fibers were constructed by the rigid and brittle PHB as the core, together with a shell layer of soft and ductile PDLLA. When the core−shell fibers are under tensile stress, the brittle core is fractured first and breaks into several segments. This phenomenon is similar to the “fragmentation test” on conventional composite tests performed to determine the adhesion between fibers and matrices. The length of PHB fragments depends on the interfacial strength between the PHB core and PDLLA shell. Based on the small aspect values of the PHB fragments, poor adhesion between the phases in core− I

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules tabulated in Table 2. Thin fibers were obtained, indicating that further stretching of the deposited fibers possibly occurs by the

direction. The chain orientation was significantly reduced for the 50/50 fibers, whereas no preferential orientation of macroscopic chains was observed in the 25/75 and 10/90 fibers because their f IR values were close to zero. Two-dimensional WAXD and SAXS were performed to characterize the orientation of unit cells and crystalline lamellae of PHB. Typical scattering patterns are displayed in Figures 12B and C, respectively. The scattering arcs are related to the PHB crystallites, because PDLLA component is in the amorphous state. On the basis of the SAXS intensity profile, long period and lamellar thickness were determined from the corresponding correlation function, and the derived values are provided in Table 2. For the neat PHB fibers, L and lc are about 5.1 and 2.9 nm, respectively. Therefore, the linear crystallinity (=lc/L) was 0.57, which is close to the fiber crystallinities obtained from DSC and FTIR, that is, ϕDSC and ϕFTIR , c c respectively, When the PDLLA content was increased to 25 wt %, the L and lc values both slightly increased to 5.3 and 3.1 nm, respectively. However, the corresponding values were reduced to 5.0 and 2.9 nm for the 50/50 core−shell fibers. For the 2D WAXD patterns, Figure S3 shows azimuthal intensity profiles of the (020), (002), and (110) reflections of PHB. For the 100/0, 90/10, 75/25, and 50/50 fibers, the (020) reflection showed peaks at 90° (transverse), and the (002) reflection showed peaks at 0° and 180° (meridian). Thus, the crystal b-axis of PHB preferentially oriented to the direction perpendicular to the fiber direction, whereas the crystal c-axis orients in the fiber axis. On the basis of these azimuthal intensity profiles, the orientation functions of f b and fc for the crystal b-axis and c-axis of PHB were calculated. For the PHB/PDLLA 25/75 fibers, the

Table 2. Fiber Diameter (dfo), Overall Chain Orientation Function (f IR), Long Period (L), and Lamellar Thickness (lc) of Aligned Fibers Collected by Rotating Wire Drum PHB/PDLLA 100/0 90/10 75/25 50/50 25/75 10/90

dfo (μm)

f IR,PLA (−)

f IR,PHB (−)

L (nm)

lc (nm)

± ± ± ± ± ±

0.10 0.17 0.03 0.04 0.04

0.22 0.24 0.18 0.07 0.01 0.00

5.1 5.2 5.3 5.0

2.9 3.0 3.1 2.9

1.48 1.72 1.69 1.94 1.76 1.65

0.22 0.22 0.29 0.24 0.28 0.26

'The italic area is for the core−shell fibers. bf IR,PLA and f IR,PHB represent the overall chain orientation of the respective PDLLA and PHB component determined from the polarized IR results. a

simple drum design. Polarized IR measurements were performed to reveal the chain orientation of two respective components in the PHB/PDLLA fibers. Figure 12A shows the typical IR polarization spectra of the 75/25, 50/50, and 25/75 fibers. After band deconvolution, the respective dichroic ratio of the PHB and PDLLA component was determined to calculate the orientation function. The calculated orientation functions for the PHB and PDLLA components, denoted by f IR,PHB and f IR,PLA, respectively, are listed in Table 2. The calculated f IR,PHB was used to represent the overall PHB chain orientation, including those in the amorphous and crystalline regions. For the 90/10 and 75/25 core−shell fibers, the PHB and PDLLA chains were both oriented preferentially parallel to the fiber

Figure 12. (A) Polarized IR spectra (solid curve: A∥; dashed curve: A⊥), (B) 2D WAXD patterns, and (C) 2D SAXS patterns of aligned PHB/ PDLLA 75/25, 50/50, and 25/75 fibers. The fiber axis is vertical. J

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules

tapering jet and eventually coalesced in the core section. By simply varying the PHB/PDLLA (w/w) ratio, the core diameter and shell thickness can be modified. The toughness of the PHB/PDLLA composite fibers significantly increased without sacrificing their elastic moduli because of their unique structure (hard-core and soft-shell structure). Core−shell fibers with tailored properties and structure for biomedical application can be obtained by appropriately controlling their compositions and solution properties (viscosity and relaxation time).

(002) reflection was extremely weak and cannot be traced. Thus, the value of ⟨cos2 ϕc⟩ was calculated from the azimuthal intensity profiles of the (020) and (110) reflections, that is, ⟨cos2 ϕc⟩ = 1−1.1902⟨cos2 ϕ110⟩ − 0.8097⟨cos2 ϕ020⟩, according to the Wilchinsky equation.31 After deriving f b and fc, the orientation function of the crystal a-axis fa was then obtained from the relationship for the orthorhombic crystal lattice of PHB:31 fa + f b + fc = 0. Figure 13 shows the calculated

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.biomac.7b00863. Status of blend solutions at room temperature. SEM images of fibers collected on the rotating wire drum. WAXD intensity profiles along the azimuthal directions of (020), (002), and (110) reflections of the α-form PHB crystal in the aligned fibers (PDF).

■

Figure 13. Orientation function of crystal a-, b-, and c-axis of the αform PHB developed in the aligned fibers with different PHB compositions.

AUTHOR INFORMATION

Corresponding Author

*Tel.: +886-6-2757575, ext. 62645. Fax: +886-6-2344496. Email: [email protected].

values of fa, f b, and fc for fibers with different PHB contents. The fc values for the PHB/PDLLA 90/10 and 75/25 fibers were higher than those for the neat PHB fibers, suggesting that the PHB/PDLLA 90/10 and 75/25 fibers exhibited better crystal orientation, and their c-axes along the fiber direction were developed. The enhanced crystal c-axis orientation is possibly associated with the effective drawing of PHB-rich phase in the core section during electrospinning, as mentioned previously. Therefore, the high Young’s moduli of the 90/10 and 75/25 core−shell fibers were attributed to their enhanced crystal orientation. When the PDLLA content (and the shell thickness) was further increased, the crystal orientation significantly changed. For the 25/75 core−shell fibers, the crystal c-axis was perpendicular to the fiber direction, and the aaxis was parallel along the fiber. This orthogonal crystal orientation was also observed in the cold drawing of the binary blends of PHB/PLA43 and PHB/cellulose acetate butyrate.27 The unique crystal orientation was confirmed from our polarized IR spectrum (Figure 12A). Clearly, the 1723 cm−1 band associated with the CO stretching mode in the crystalline PHB was larger for the A∥ than that for the A⊥. As shown in Figure 9, the two-stage drawing phenomenon observed in the 25/75 core−shell fibers is possibly associated with the reorientation of PHB crystal during tensile stretching.

ORCID

Chi Wang: 0000-0003-0627-2341 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors are grateful to the Ministry of Science and Technology, Taiwan, R.O.C. for the research grant (MOST103-2221-E-006-262-MY3) that supported this work. For the 2D SAXS/WAXD experiments, the support from National Synchrotron Radiation Research Center and helpful assistance from Drs. U-Ser Jeng, Chun-Jen Su, and Jey-Jau Lee are highly appreciated.

■

REFERENCES

(1) Agarwal, S.; Wendorff, J. H.; Greiner, A. Use of Electrospinning Technique for Biomedical Applications. Polymer 2008, 49, 5603− 5621. (2) Xu, C. Y.; Inai, R.; Kotaki, M.; Ramakrishna, S. Aligned Biodegradable Nanotibrous Structure: a Potential Scaffold for Blood Vessel Engineering. Biomaterials 2004, 25, 877−886. (3) Zussman, E.; Yarin, A. L.; Bazilevsky, A. V.; Avrahami, R.; Feldman, M. Electrospun Polyaniline/Poly(methyl methacrylate)Derived Turbostratic Carbon Micro-/Nanotubes. Adv. Mater. 2006, 18, 348−353. (4) Wang, M.; Yu, J. H.; Kaplan, D. L.; Rutledge, G. C. Production of Submicron Diameter Silk Fibers under Benign Processing Conditions by Two-Fluid Electrospinning. Macromolecules 2006, 39, 1102−1107. (5) Ojha, S. S.; Stevens, D. R.; Hoffman, T. J.; Stano, K.; Klossner, R.; Scott, M. C.; Krause, W.; Clarke, L. I.; Gorga, R. E. Fabrication and Characterization of Electrospun Chitosan Nanofibers Formed via Templating with Polyethylene Oxide. Biomacromolecules 2008, 9, 2523−2529. (6) Moghe, A. K.; Gupta, B. S. Co-axial Electrospinning for Nanofiber Structures: Preparation and Applications. Polym. Rev. 2008, 48, 353−377. (7) Jiang, H.; Hu, Y.; Li, Y.; Zhao, P.; Zhu, K.; Chen, W. A Facile Technique to Prepare Biodegradable Coaxial Electrospun Nanofibers

4. CONCLUSIONS Scaffolds made of core−shell electrospun fibers are currently attracting interest in the fields of tissue engineering and cell biology. Core−shell fibers were fabricated by performing singletube electrospinning on phase-separated PHB/PDLLA solutions. For the blend solution with the PHB content of ≥25 wt %, core−shell fibers with crystalline PHB cores and amorphous PDLLA shells were constantly obtained. The formation of core−shell structure was possibly caused by PHB gelation in the PHB-rich phase subjected to a high extension rate during electrospinning. This phase formed string-like structures in the jet. The string-like structures were surrounded by the less viscous PDLLA-rich phase and tended to migrate inward in the K

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX

Article

Biomacromolecules for Controlled Release of Bioactive Agents. J. Controlled Release 2005, 108, 237−243. (8) Chew, S. Y.; Wen, J.; Yim, E. K. F.; Leong, K. W. Sustained Release of Proteins from Electrospun Biodegradable Fibers. Biomacromolecules 2005, 6, 2017−2024. (9) Huang, Z. M.; He, C. L.; Yang, A.; Zhang, Y.; Han, X. J.; Yin, J.; Wu, Q. Encapsulating Drugs in Biodegradable Ultrafine Fibers Through Co-axial Electrospinning. J. Biomed. Mater. Res., Part A 2006, 77A, 169−179. (10) Zhang, Y. Z.; Wang, X.; Feng, Y.; Li, J.; Lim, C. T.; Ramakrishna, S. Coaxial Electrospinning of (Fluorescein Isothiocyanate-Conjugated Bovine Serum Albumin)-Encapsulated Poly(εcaprolactone) Nanofibers for Sustained Release. Biomacromolecules 2006, 7, 1049−1057. (11) Chakraborty, S.; Liao, I. C.; Adler, A.; Leong, K. W. Electrohydrodynamics: A Facile Technique to Fabricate Drug Delivery Systems. Adv. Drug Delivery Rev. 2009, 61, 1043−1054. (12) Wei, M.; Kang, B.; Sung, C.; Mead, J. Core-Sheath Structure in Electrospun Nanofibers from Polymer Blends. Macromol. Mater. Eng. 2006, 291, 1307−1314. (13) Xu, X.; Zhuang, X.; Chen, X.; Wang, X.; Yang, L.; Jing, X. Preparation of Core-Sheath Composite Nanofibers by Emulsion Electrospinning. Macromol. Rapid Commun. 2006, 27, 1637−1642. (14) Bazilevsky, A. V.; Yarin, A. L.; Megaridis, C. M. Coelectrospinning of Core-Shell Fibers Using a Single-Nozzle Technique. Langmuir 2007, 23, 2311−2314. (15) Chen, M.; Dong, M.; Havelund, R.; Regina, V. R.; Meyer, R. L.; Besenbacher, F.; Kingshott, P. Thermo-Responsive Core-Sheath Electrospun Nanofibers from Poly(N-isopropylacrylamide)/Polycaprolactone Blends. Chem. Mater. 2010, 22, 4214−4221. (16) Ishii, D.; Lee, W. K.; Kasuya, K. I.; Iwata, T. Fine Structure and Enzymatic Degradation of Poly[(R)-3-hydroxybutyrate] and Stereocomplexed Poly(lactide) Nanofibers. J. Biotechnol. 2007, 132, 318− 324. (17) Zong, X.; Kim, K.; Fang, D.; Ran, S.; Hsiao, B. S.; Chu, B. Structure and Process Relationship of Electrospun Bioabsorbable Nanofiber Membranes. Polymer 2002, 43, 4403−4412. (18) Inai, R.; Kotaki, M.; Ramakrishna, S. Structure and Properties of Electrospun PLLA Single Nanofibres. Nanotechnology 2005, 16, 208− 213. (19) Wang, C.; Hsu, C. H.; Hwang, I. H. Scaling Laws and Internal Structure for Characterizing Electrospun Poly[(R)-3-hydroxybutyrate] Fibers. Polymer 2008, 49, 4188−4195. (20) Wang, C.; Chien, H. S.; Yan, K. W.; Hung, C. L.; Hung, K. L.; Tsai, S. J.; Jhang, H. J. Correlation Between Processing Parameters and Microstructure of Electrospun Poly(D,l-lactic acid) Nanofibers. Polymer 2009, 50, 6100−6110. (21) Wang, C.; Yan, K. W.; Lin, Y. D.; Hsieh, C. H. Biodegradable Core/Shell Fibers by Coaxial Electrospinning: Processing, Fiber Characterization, and Its Application in Sustained Drug Release. Macromolecules 2010, 43, 6389−6397. (22) Katta, P.; Alessandro, M.; Ramsier, R. D.; Chase, G. G. Continuous Electrospinning of Aligned Polymer Nanofibers onto a Wire Drum Collector. Nano Lett. 2004, 4, 2215−2218. (23) Barham, P. J.; Keller, A.; Otun, E. L.; Holmes, P. A. Crystallization and Morphology of a Bacterial Thermoplastic: Poly3-hydroxybutyrate. J. Mater. Sci. 1984, 19, 2781−2794. (24) Zong, X.; Ran, S.; Kim, K. S.; Fang, D.; Hsiao, B. S.; Chu, B. Structure and Morphology Changes during in Vitro Degradation of Electrospun Poly(glycolide-co-lactide) Nanofiber Membrane. Biomacromolecules 2003, 4, 416−423. (25) Pan, P.; Kai, W.; Zhu, B.; Dong, T.; Inoue, Y. Polymorphous Crystallization and Multiple Melting Behavior of Poly(l-lactide): Molecular Weight Dependence. Macromolecules 2007, 40, 6898−6905. (26) Yoshie, N.; Oike, Y.; Kasuya, K.; Doi, Y.; Inoue, Y. Change of Surface Structure of Poly(3-hydroxybutyrate) Film upon Enzymatic Hydrolysis by PHB Depolymerase. Biomacromolecules 2002, 3, 1320− 1326.

(27) Park, J. W.; Tanaka, T.; Doi, Y.; Iwata, T. Uniaxial Drawing of Poly[(R)-3-hydroxybutyrate]/Cellulose Acetate Butyrate Blends and Yjeir Orientation Behabior. Macromol. Biosci. 2005, 5, 840−852. (28) Jeng, U. S.; Su, C. H.; Su, C. J.; Liao, K. F.; Chuang, W. T.; Lai, Y. H.; Chang, J. W.; Chen, U. J.; Huang, Y. S.; Lee, M. T.; Yu, K. L.; Lin, J. M.; Liu, D. G.; Chang, C. F.; Liu, C. Y.; Chang, C. H.; Liang, K. S. A Small/wide-angle X-ray Scattering Instrument for Structural Characterization of Air−liquid Interfaces, Thin Films and Bulk Specimens. J. Appl. Crystallogr. 2010, 43, 110−121. (29) Strobl, G. R.; Schneider, M. Direct evaluation of the electron density correlation function of partially crystalline polymers. J. Polym. Sci., Polym. Phys. Ed. 1980, 18, 1343−1359. (30) Wang, C.; Hsieh, T. C.; Cheng, Y. W. Solution-Electrospun Isotactic Polypropylene Fibers: Processing and Microstructure Development during Stepwise Annealing. Macromolecules 2010, 43, 9022−9029. (31) Alexander, L. E. X-ray Diffraction Methods in Polymer Science; John Wiley & Sons, Inc.: New York, 1969; p 241. (32) Koyama, N.; Doi, Y. Miscibility of Binary Blends of Poly[(R)-3hydroxybutyric acid] and Poly[(S)-lactic acid]. Polymer 1997, 38, 1589−1593. (33) Turchetto, A.; Cesaro, A. Gel-sol Phase Transition of Poly(3hydroxybutyrate) in Dimethylformamide. Thermochim. Acta 1995, 269, 307−317. (34) Bognitzki, M.; Frese, T.; Steinhart, M.; Greiner, A.; Wendorff, J. H. Preparation of Fibers with Nanoscaled Morphologies: Electrospinning of Polymer Blends. Polym. Eng. Sci. 2001, 41, 982−989. (35) Sun, X.; Tokuda, A.; Oji, Y.; Nakatani, T.; Tsuji, H.; Ozaki, Y.; Yan, S.; Takahashi, I. Effects of Molar Mass of Poly(L-lactide acid) on the Crystallization of Poly[(R)-3-hydroxybutyrate] in Their Ultrathin Films. Macromolecules 2012, 45, 2485−2493. (36) He, J. H.; Wu, Y.; Zuo, W. W. Critical Length of Straight Jet in Electrospinning. Polymer 2005, 46, 12637−12640. (37) Williams, M. C. Migration of 2 Liquid-phases in Capillary Extrusion-Energy Interpretation. AIChE J. 1975, 21, 1204−1207. (38) Kim, C.; Jeong, Y. I.; Ngoc, B. T. N.; Yang, K. S.; Kojima, M.; Kim, Y. A.; Endo, M.; Lee, J. W. Synthesis and Characterization of Porous Carbon Nanofibers with Hollow Cores Through the Thermal Treatment of Electrospun Copolymeric Nanofiber Webs. Small 2007, 3, 91−95. (39) Hashimoto, T.; Matsuzaka, K.; Moses, E.; Onuki, A. String Phase in Phase-Separating Fluids under Shear Flow. Phys. Rev. Lett. 1995, 74, 126−129. (40) Wang, C.; Wang, Y.; Hashimoto, T. Impact of Entanglement Density on Solution Electrospinning: a Phenomenological Model for Fiber Diameter. Macromolecules 2016, 49, 7985−7996. (41) Yokouchi, M.; Chatani, Y.; Tadokoro, H.; Teranishi, K.; Tani, H. Structural Studies of Polyesters: 5. Molecular and Crystal Structures of Optically Active and Racemic Poly (β-hydroxybutyrate). Polymer 1973, 14, 267−272. (42) Orts, W. J.; Marchessault, R. H.; Bluhm, T. L.; Hamer, G. K. Observation of strain-induced β form in poly(β-hydroxyalkanoates). Macromolecules 1990, 23, 5368−5370. (43) Vogel, C.; Wessel, E.; Siesler, H. W. FT-IR Spectroscopic Imaging of Anisotropic Poly(3-hydroxyrate)/Poly(lactic acid) Blends with Polarized Radiation. Macromolecules 2008, 41, 2975−2977.

L

DOI: 10.1021/acs.biomac.7b00863 Biomacromolecules XXXX, XXX, XXX−XXX