Electrospinning of Poly(L-lactide-co-DL-lactide) Copolymers: Effect of Chemical Structures and Spinning Conditions.
Polylactide (PLA) is a thermoplastic aliphatic polyester, which is widely used in packaging, medical, and agricultural applications, due to its degradability, biocom-patibility, and renewability (1-3). PLA can be synthesized by polycondensation of lactic acid, derived from starch fermentation, or ring-opening polymerization of its cyclic lactide (LA) dimer, where the latter approach is common for producing high molecular weight (MW) products. LA is present in three forms, that is, L-lactide (LLA), D-lactide (DLA), and DL-lactide (DLLA). The chain structure and properties of PLA can be optimized by employing different LA forms (4). Because of the LLA abundance, poly(L-lactide), (PLLA), is widely used in commercial applications. Although the polymer has good mechanical properties, comparable to other commodity plastics (5), it is too rigid and brittle to be used in certain applications (6), Copolymerization of LLA and DLLA at various compositions produces poly(L-lactide-co-DL-Lactide) (PDLL[A.sub.x]), which possesses adjustable crystallinity and mechanical properties (7). The copolymers have high potential for biomedical applications, such as cell scaffold, drugs controlled-release materials, and membrane filters, especially those fabricated in the form of nanofibers.
Electrospinning is recognized as an efficient technique for fabricating nanofibers from various polymers, either in the forms of solution or melt. Polymeric nanofibers show superior characteristics, that is, very large surface area to volume ratio, high flexibility, and high mechanical performance, compared to traditional materials, which make the materials well-known for many applications (8), (9). Morphology and properties of electrospun fibers can be optimized by adjusting the processing parameters, such as solvent system and concentration, voltage, and the collecting process of the spun fibers (10). Electrospun fibers of biodegradable/biocompatible polymers, especially PLA, have high potential for use in biomedical applications (11), (12). Electrospinning of PLLA homopolymer has been studied by employing various solvents, for example, DMF (13), THF (14), DMF/THF (15), and DMF/CH[Cl.sub.3] (16).
In this study, nanofibers of PDLL[A.sub.x] copolymers are fabricated by electrospinning. Commercial PLLA and in-house synthesized PDLL[A.sub.x] copolymers, consisting of various DLLA contents are employed. Effects of (co)polymer nature (MW, DLLA contents, crystallizability) and electrospinning parameters, that is, solvent systems and humidity, on size and surface morphology of the nanofibers are investigated. Tensile behaviors, thermal properties, chain orientation, and crystalline characteristics of the resulting nanofibers are examined.
MATERIALS AND METHODS
Commercial PLLA 4042D ([[bar.M].sub.n] = 120,000), cPLLA, was purchased from NatureWork. Medium-sized poly(t.- lactide), mPLLA, ([[bar.M].sub.n] = 40,000) and PDLLA,, copolymers with 2.5, 7.5, and 50% DLLA contents (denoted as PDLLA2.5, PDLLA7.5, and PDLLA50) were synthesized by the method previously described (3), (7). Chloroform (CH[Cl.sub.3]) and dimethyl formamide (DMF) solvents were purchased from Carlo Erba.
Fabrication of Nanofibers
Electrospun nanofibers were fabricated by placing polymer solutions in a syringe, which was placed on a syringe pump (Kd Scientific). The critical concentration ([C.sub.crit]), that is, the minimum concentration suitable for spinning due to optimum degree of chain entanglement (15), (17), of polymer solution in each solvent system was determined, and employed as a minimum processing concentration. A flow rate of I ml/h and the electrical potential of 5-7 kV were applied to a hypodermic needle on the end of the syringe using a high voltage power supply (GAMMA High Voltage Research). A grounded aluminum plate collector was placed 10-12 cm away from the needle tip. Nanofibers were deposited on an aluminum foil and collected in the form of nonwoven mats. The electrospinning process was conducted in a vented box at 25[degrees]C with an ambient relative humidity (RH) of ~40-50%. The resulting spun fibers were then dried in a vacuum oven for 24 h to remove the residual solvent. In the study of effect of RH on properties of the nanofibers, RH was varied from 30 to 80% by using a humidifier placed inside the electrospinning vent box.
Morphology of electrospun fiber mats was examined by SEM (HITACHI S3400N). The samples were prepared by a platinum coating technique. Average size and size distribution were determined by using the Image Analyzer (PC SEM) program. Surface porosity was examined by an image processing technique using the Adobe Photoshop program. An area of interest was first selected, and the contrast of the image was then adjusted to distinguish smooth surface and pore areas. The number of pixels of the two patterns was determined, and the surface porosity was finally calculated.
Crystalline characteristics of the spun fiber samples were examined by XRD on a JEOL JDX-2520 diffractometer, using CuK21 radiation. The (co)polymer samples in the forms of solution-cast films and as-spun fiber mats were scanned from 2[theta] of 5-50[degrees] with a 0.02[degrees] step size. Thermal properties were examined by DSC on a Mettler Toledo [DSC822.sup.e]. The samples were scanned twice from -20 to 200[degrees]C with a heating and cooling rate of 10[degrees]C/min. Tensile measurements were carried out on an Instron tensile testing machine (Model 55R4502), according to ASTM D882, with a crosshead speed of 50 mm/min, using rectangle specimens with 50 mm gauge length and 15 mm width.
RESULTS AND DISCUSSION
Nanofibers of PLLAs with Different Molecular Weights
Effects of PLLA's MW and solvent systems on properties of the spun fibers were investigated. cPLLA and mPLLA were electrospun using two solvent systems, that is, CH[Cl.sub.3] and a 1:3 (w/w) DMF/CHC13 mixed solvent, with different evaporation rates. The chloroform boiling point, Hildebrand solubility, and vapor pressure are 61[degrees]C, 19.0 [(MPa).sup.1/2], and 21.2 KPa, respectively, and for DMF, 153[degrees]C, 24.7 [(MPa).sup.1/2], and 0.3 KPa, respectively. Bead defects and fiber nonuniformity are apparent for fibers of mPLLA derived from both solvents, as shown in Fig. I. The result reflects that the polymer chain length plays a key role in the nanofiber's properties. As the MW of mPLLA is not high enough to form tight-chain entanglement, a heterogeneous jet is formed, leading to bead formation. For high-MW cPLLA, defect-free fibers were obtained from both solvents. The mixed solvent produces relatively smooth fibers that are ~10 times smaller in average size, compared to the porous counterparts derived from CH[Cl.sub.3]. During the electrospinning, elongational flow of non-Newtonian solutions retards the breakup of their viscoelastic jets, leading to formation of long threads of jets. This phenomenon is strongly dependent on polymer concentration and MW (18).
Nanofibers of PDLL[A.sub.x] Copolymers
PDLL[A.sub.x] copolymers possessed intriguing structure/properties correlation, where tensile behaviors were strongly dependent on DLLA content (7). The copolymers consisting of 2.5 and 7.5% DLLA contents showed an increase in tensile modulus, compared to mPLLA. PDLLA50 exhibited the maximum elongation at break in this copolymer series. These three copolymers with comparable MW, ranging from 80 to 118 X [10.sup.3] g/mol, were fabricated into nanofiber mats. Table 1 summarizes effects of the copolymer's structure (DLLA content) and solvent systems on size, morphology, and properties of the spun fibers. Bead defect-free fibers were obtained from all copolymers regardless of the solvent systems, as their MW is high enough to provide a sufficient degree of chain entanglements. The average size of nanofibers derived from CH[Cl.sub.3](6.3-7.5 [micro]m), however, is roughly an order of magnitude larger than that from DMF/CH[Cl.sub.3]. This reflects a profound effect of the solvent evaporation and jet solidification rates on the fiber's properties. It was reported that the size of spun polymeric fibers decreases with the use of a solvent with a slow evaporation rate, due to higher vapor concentration of solvent (19). Thermally induced and vapor induced phase separation mechanisms were proposed to explain the critical effect of solvent evaporation rate on the fiber size and morphology (20), (21).
TABLE 1. Electrospinning conditions and properties of the resulting nanotibers. Samples [[bar.M].sub.n] Solvent Cone. Fiber (wt%) appearance mPLLA 40000 CH[CL.sub.3] 13 Some beads DMF/CH[CL.sub.3] 15 Some beads PDLLA2.5 83340 CH[CL.sub.3] 8 No bead DMF/CH[CL.sub.3] 8 No bead PDLLA7.5 118250 CH[CL.sub.3] 10 No bead DMF/CH[CL.sub.3] 18 No bead PDLLA50 80880 CH[CL.sub.3] 15 No bead DMF/CH[CL.sub.3] 15 No bead Samples Fiber diameter (nm) mPLLA 6800 [+ or -] 3000 747 [+ or -] 230 PDLLA2.5 7500 [+ or -] 1200 723 [+ or -] 200 PDLLA7.5 7000 [+ or -] 1600 854 [+ or -] 430 PDLLA50 6300 [+ or -] 1200 568 [+ or -] 130
Figure 2 shows SEM images of (bead-free) spun fibers of PDLL[A.sub.x], copolymers. Solvent systems impose stronger influence on the fiber's surface features, compared to the copolymer's nature. CH[Cl.sub.3] produces fibers with a high degree of porosity, whereas those derived from the mixed solvent exhibit a significantly smooth surface. This is due to a different phase separation mechanism, which is a result from different solvent evaporation rates (20). The copolymer structures partly affect the fiber properties when the same solvent system is employed, especially with the use of CH[Cl.sub.3], as a high degree of property variations is observed. The spun fibers of semicrystalline copolymers, that is, PDLLA2.5 and PDLLA7.5, exhibit high surface porosity, compared to an amorphous PDLLA50. This indicates that surface porosity is also dependent on crystallize-ability of the copolymers, when a solvent with a fast evaporation rate is employed. The explanation involves a competition between polymer crystallization and solvent evaporation/solidification of the solution jets. It is noted that the surface pores developed here are also affected by the ambient humidity, which influences the phase separation mechanisms, as discussed below.
Effect of Humidity
A decrease in fiber diameter and the presence of surface pores of spun nanofibers leads to a large increase in their surface area, which directly enhances their application use. This is adjustable by varying the RH conditions of the electrospinning. Effect of humidity on properties of PDLL[A.sub.x], nanofibers using CH[Cl.sub.3] was examined by varying RH values from 30, 50, to 80%. The results are summarized in Table 2. The fiber diameters of crystallizable mPLLA and PDLLA2.5 increase with RH values, whereas that of an amorphous PDLLA50 shows a reverse trend. This agrees with a recent study that the change in RH condition can change the nanofiber thickness, depending on the chemical nature of the polymer, polymer-solvent interaction, which leads to different evaporation rates (22). The increase in RH typically causes an increase in fiber size and formation of surface features of spun fibers derived from nonaqueous solvent, for example, poly(acrylonitrile)/DMF, polysulfone/DMF (23), and polystyrene/ THF (24). The interplay between crystallization and discharging rate of the solution jets was proposed as the origin of this behavior. During the spinning, the applied surface charge on solution jets is interfered with and captured by moisture droplets in the surroundings, as water has higher conductivity than air. This lowers the opportunity to split the solution jets to become thinner fiber, leading to formation of larger fibers, for semicrystalline copolymers. (24) In addition, it was reported that RH conditions affected fiber properties differently, depending on the polymer structure and interactions, due to complex interactions between the nonsolvent (water), the hygroscopic solvent, and the polymer. For example, polysulfone was affected stronger by the RH than poly(acrylonitrile) (23).
TABLE 2. Electrospinning conditions and properties of spun fibers prepared by using CH[Cl.sub.3], as a function of relative humidity (RH) and copolymer structure. Sample Cone. RH (%) Diameter Porosity (%) (wt%) ([micro]m) mPLLA 13 30 2.4 [+ or -] 0 0.8 50 4.8 [+ or -] 39 1.3 80 5.5 [+ or -] 45 1.6 PDLLA2.5 8 30 4.3 [+ or -] 2 1.4 50 4.6 [+ or -] 19 2.6 80 5.0 [+ or -] 43 0.9 PDLLA50 15 30 6.5 [+ or -] 0 0.4 50 6.3 [+ or -] 16 1.6 80 4.8 [+ or -] 36 2.1
In a fabrication of hydrophilic polymers, for example, poly(ethylene oxide) using water as a solvent, a decrease in the fiber size with increasing RH was observed. This is from a decrease in the evaporation rate of the jet's water solvent, but there was no effect of solvent/nonsolvent interplay (19). The decrease in fiber diameter with increasing RH of amorphous PDLLA50 is likely explained by a similar mechanism. The lack of crystallize-ability of the copolymer leads to a domination of solvent evaporation in the solidification of the solution jets, allowing the jets to spin at greater discharge conditions to form smaller fibers. The increase in RH results in an even slower solvent evaporation rate, which allows the fibers to further elongate and to be thinner.
Figure 3 and Table 2 show the critical effect of humidity on surface features of the copolymer fibers fabricated from the fast evaporating CH[Cl.sub.3] solvent. Surface porosity of all samples increases with increasing RH. A rapid increase in porosity is observed when RH was changed from 30 to 50%, whereas a slight increase is obtained from RH of 50-80%, due to a saturation of pore content on the fiber surface. Vapor-induced phase separation of the homogeneous PDLLA jets involving the penetration of a nonsolvent water vapor and a rapid evaporation of the solvent from the jets, which causes a drop in temperature of the traveling jets, that is, thermally induced phase separation, have been proposed as the origins of the surface features formation (21).
Thermal Properties and Chain Structure
Thermal properties of the spun fibers of PDLL[A.sub.x] with different DLLA contents were examined in comparison with the corresponding solution-cast film samples, where the DSC thermograms (first heating scan) are shown in Fig. 4. Melting and glass transition characteristics are observed in all semicrystalline samples. A decrease of glass transition temperature ([T.sub.g]) and melting temperature ([T.sub.m]) with DLLA content of both fibers and film samples was found, as summarized in Table 3. The spun fibers have higher [T.sub.g] and lower crystalline content than the as-cast film samples. This is because the copolymer chains are largely elongated by the applied electric field along the solution jets during the spinning process. Upon rapid solvent evaporation, the chains are quenched, leading to restriction of chain's segmental movements and chain-folding arrangement to form crystalline lamellae (25-28).
All DSC thermograms of as-spun fibers show an endothermic aging peak at a temperature slightly higher than their [T.sub.g]. This is not observed in those of the corresponding film counterparts (first heating scan) and the second heating run thermograms of spun fibers (not shown). This is due to relaxation of the strain imposed from the rapid solidification of the elongated copolymer chains (29). Residue charges trapped within the fibers, probably at amorphous/crystalline interface, may also be responsible for this characteristic, as these are released near the glass transition due to a dipole relaxation process (30). It is noted that the enthalpy heat of this relaxation peak, as shown in Table 3, is significantly correlated with the crystallinity difference of the film and fibers samples, reflecting the influence of the copolymer nature. When the as-spun fibers are heated up above their glass transition and the aging endothermic relaxation processes, reorientation is possible, and a cold crystallization of the samples takes place. This is not observed in those of the cast film counterparts, as the electrospinning process suppresses chain-folding arrangement to form crystal lamellae of the asspun fibers (25), (26). Improvement of crystalline content was achieved by annealing of the samples (11), (27).
XRD traces of as-spun fiber mats prepared from CH[Cl.sub.3] and the mixed solvent, as a function of DLLA content, are compared in Figs. 5 and 6. These are different from that of PLLA crystallite ([10.sub.3] helix structure), which exhibits peaks at 2[theta] of 15, 17, and 19[degrees], whereas a [3.sub.t] helical conformation of PDLA/PLLA stereocomplex shows a characteristic pattern at 12, 21, and 24[degrees]. (31) The as-spun fibers show a peak pattern, with very weak intensity, at 2[theta] of 15.2, 17.2, and l9.3[degrees]. This reflects the presence of PLLA crystallites with low crystalline content, which agrees with DSC results. However, very sharp signals are observed at 2[theta] of 38.2 and 44.4[degrees], corresponding to dspacing of 2.35 and 2.03 nm, respectively. These unique signals are likely derived from an orientation of intramolecular hydrogen bonding between methine --C--H and C=0 groups in the copolymer repeat units, as shown in Fig. 7. The formation of this conformation is likely due to the rapid solidification of elongated chains as a result from the applied high electric field. This, in turn, imposes conformational strain on the copolymer chains to resume a preferred random coil conformation, which retards their reorientation to form crystal lamellae, as previously discussed. The intensity of these signals is strongly dependent on the fabricated solvent system, where a 10-fold sharper signal is observed when CH[Cl.sub.3] is employed, as its fast evaporation rate imposes a higher degree of quenching. Consequently, longer intramolecular hydrogen bonding sequences are generated.
Mechanical properties of PDLL[A.sub.x] copolymers in the forms of cast film and spun fiber mats are characterized and summarized in Table 4. Enhancement in bulk modulus, strength, and elongation at break, with increasing DLLA contents, is obtained in semicrystalline copolymer films. This may be due to the combined effects of crystalline domains and "physical crosslinks" formed between DLLA sequences in the chains. The DL lactate segments and their mirror image LD lactate counterparts are capable of forming tight-chain interaction by strong intermolecular hydrogen bonding. These "physical crosslinks" are largely present in amorphous domains, whose contents are dependent on the DLLA composition in the copolymer chains. For PDLLA50, however, a drop in the modulus and strength is observed because of its lack of crystalline domains, whereas a much improved elongation at break is achieved, possibly due to high content and isotropic distribution of its "physical crosslink." Figure 8a shows tensile behaviors of PDLLA7.5 in the forms of cast film and fiber mats, where the former exhibits (an order of magnitude) higher tensile stress. A sharp yield point is observed in the film sample, associated with a break of crystalline domains due to its semi-crystalline nature. The stress-strain curves of spun fibers are compared in Fig. 8b. No sharp yield point is observed, reflecting an influence of a structural factor. Different orientation of the polymer chains and crystalline content of the samples in both forms are partly contributing to this behavior (32).
TABLE 3. Thermal properties of PDLL[A.sub.x] copolyrners in the forms of as-cast film and as-spun fiber mats fabricated from CH[Cl.sub.3] solvent, determined from first heating scan DSC thermograms. Film Fiber mats [T.sub.g] [T.sub.m] [X.sub.c] (a) Relaxation Sample ([degrees]C) ([degrees]C) (%) heat (J/g) mPLLA 68 178 42 4.1 PDLLA2.5 61 168 40 3.5 PDLLA7.5 58 154 30 7.0 PDLLA50 54 - 0 6.9 [T.sub.g] [T.sub.m] [X.sub.c] (a) Sample ([degrees]C) ([degrees]C) (%) mPLLA 64 179 22 PDLLA2.5 65 168 26 PDLLA7.5 62 159 8 PDLLA50 57 - 0 (a) Crystallinity, [X.sub.c] = ([DELTA][H.sub.m] - [DELTA][H.sub.c])/ [DELTA][H.sub.m], [DELTA][H.sub.m] = 93.6 J/g.
Tensile characteristics of fiber mats are summarized in Table 4. The magnitudes of strength and modulus are much lower, compared to the polymer films, because the fibers are loosely packed and interact only one-dimensionally (33). Elongation at break is comparable to or higher than that of the cast films. Effect of the copolymer structure, that is, the DLLA content, on mechanical properties is different from that of the film samples, due to the competing effect of the fiber structure and interplay between intra-/ intermolecular hydrogen bonding. PDLLA50 fiber mats show an enhancement in all mechanical properties, whose modulus, strength, and elongation at break are superior to other fiber samples. Their high DL-lactate content might be responsible for forming of more "physical crosslinks." Detailed information on this phenomenon requires further investigation. Nevertheless, these fiber mats of amorphous copolymer with high mechanical properties have a high potential of use for many biomedical applications.
TABLE 4. Mechanical properties of PDLL[A.sub.X] copolymers in the forms of as-cast film and fiber mats, as a function of DLLA contents. Modulus (MPa) Strength (MPa) Sample Film Fiber mats Film PDLLA2.5 2500 [+ or -] 8 [+ or -] 1 46.5 [+ or -] 10 6.8 PDLLA7.5 3370 [+ or -] 43 [+ or -] 2 62.7 [+ or -] 70 2.2 PDLLA50 1550 [+ or -] 140 [+ or -] 10 22.4 [+ or -] 40 5.1 Elongation at break (%) Sample Fiber mats Film Fiber mats PDLLA2.5 0.3 [+ or -] 15 [+ or -] 5 20 [+ or -] 3 0.1 PDLLA7.5 0.6 [+ or -]0.3 30 [+ or -] 6 18 [+ or -] 4 PDLLA50 1.7 [+ or -] 148 [+ or -] 5 113 [+ or -] 9 0.4
Nanofibers of PDLL[A.sub.x] copolymers are fabricated by electrospinning. Effects of copolymer structure and its molecular weight, solvents, and humidity on morphology and properties of the fibers are investigated. The spun fibers of mPLLA with MW lower than 40,000 g/mol show bead defects, regardless of the solvent type, as the required tight-chain entanglements cannot be achieved. Defect-free fibers of PDLL[A.sub.x] fabricated using a DMF:CH[Cl.sub.3]mixed solvent are roughly 10-times smaller in diameters, with lower degree of surface porosity, compared to those of CH[Cl.sub.3] Relative humidity plays a major role on average size and surface features of the spun fibers, which strongly determine their potential applications. An increase in size and surface porosity with RH value is observed in crystallizable copolymers, whereas that of the amorphous copolymer shows a reverse trend. Thermal properties and chain arrangement of the spun fibers are critically affected by DLLA content of the copolymers and electrospinning conditions, where interplay between intermolecular and intramolecular hydrogen bonding plays a major role. PDLLA50 spun fibers show a large improvement in all aspect of mechanical properties, which is suitable for biomedical applications, pertaining to its amorphous nature and good mechanical properties.
Correspondence to: Pakorn Opaprakasit; e-mail: firstname.lastname@example.org Contract grant sponsor: National Research University Project of Thailand, Office of Higher Education Commission; contract grant sponsor: Thailand Research Fund (TRF)/Thailand Office of Higher Education Commission; contract grant number: RTA5480007; contract grant sponsor: SIIT, Thammasat University.
DOI 10.1002/pen .23576
Published online in Wiley Online Library (wileyonlinelibrary.com).
[c] 2013 Society of Plastics Engineers
(1.) K. Madhavan Nampoothiri, N.R. Nair, and R.P. John, Bioresour. Technol., 101, 8493 (2010).
(2.) A.G. Andreopoulos, E. Hatzi, and M. Doxastakis, J. Mater. Sci Mater. Med., 10, 29 (1999).
(3.) R.E. Drumright, P.R. Gruber, and D.E. Henton, Adv. Mater., 12, 1841 (2000).
(4.) A.J. Amass, K.L.R. N'goala, B.J. Tighe, and F. Schue, Polymer, 40, 5073 (1999).
(5.) R. Mehta, V. Kumar, H. Bhunia, and S.N. Upadhyay, Polym. Rev., 45, 325 (2005).
(6.) X.M. Deng, C.D. Xiong, L.M. Cheng, H.H. Huang, and R.P. Xu, J. Appl. Polym. Sc., 55, 1193 (1995).
(7.) S. Buchatip, A. Petchsuk, and K. Kongsuwan, J. Met. Mater. Miner., 18, 175 (2008).
(8.) P.D. Dalton, D. Grafahrend, K. Klinkhammer, D. Klee, and M. Moller, Polymer, 48, 6823 (2007).
(9.) S.A. Theron, E. Zussman, and A.L. Yarin, Polymer, 45, 2017 (2004).
(10.) T. Subbiah, G.S. Bhat, R.W. Tock, S. Parameswaran, and S.S. Ramkumar, J. Appl. Polym. Sci., 96, 557 (2005).
(11.) M. Spasova, N. Manolova, D. Paneva, R. Mincheva, P. Dubois, 1. Rashkov, V. Maximova, and D. Danchev, Biomacromolecules, 11, 151 (2010).
(12.) A. Toncheva, M. Spasova, D. Paneva, N. Manolova, and 1. Rashkov, J. Bioact. Compat. Polym., 26, 161 (2011).
(13.) M.W. Frey, C. Xiang, M.P. Hoffmann, A.G. Taylor, and J. Gardner, United States Patent (2011).
(14.) A. Greiner, and J.H. Wendorff, Angew. Chem. int. Ed. Engl., 46, 5670 (2007).
(15.) G.-T. Kim, J.-S. Lee, J.-H. Shin, Y.-C. Ahn, Y.-J. Hwang, H.-S. Shin, J.-K. Lee, and C.-M. Sung, Korean J. Chem. Eng., 22, 783 (2005).
(16.) D. Garlotta, J. Polym. Environ., 9, 63 (2001).
(17.) C. Zhang, X. Yuan, L. Wu, Y. Han, and J. Sheng, Eur. Polym. J., 41, 423 (2005).
(18.) R.P. Mun, J.A. Byars, and D.V. Boger, J. Non-Newtonian Fluid Mech., 74, 285 (1998).
(19.) S. Tripatanasuwan, Z. Zhong, and D.H. Reneker, Polymer, 48, 5742 (2007).
(20.) S. Megelski, J.S. Stephens, D. Bruce Chase, and J.F. Rabolt, Macromolecules, 35, 8456 (2002).
(21.) J.T. Mccann, M. Marquez, and Y. Xia, J. Am. Chem. Soc., 128, 1436 (2006).
(22.) S. Vrieze, T. Camp, A. Nelvig, B. Hagstrom, P. Westbroek, and K. Clerck, J. Mater. Sci., 44, 1357 (2008).
(23.) L. Huang, N.-N. Bui, S.S. Manickam, and J.R. Mccutcheon, J. Polym. Sci. Part B: Polym. Phys., 49, 1734 (2011).
(24.) G.T. Kim, J.S. Lee, J.H. Shin, Y.C. Ahn, K.H. Jeong, C.M. Sung, and J.K. Lee, Microsc. Microanal., 10, 554 (2004).
(25.) Z.M. Huang, Y.Z. Zhang, M. Kotaki, and S. Ramakrishna, Compos. Sci. Technol., 63, 2223 (2003).
(26.) A. Greiner and J.H. Wendorff, Angew. Chem. int. Ed. Engl., 46, 5670 (2007).
(27.) A.-R. Cho, D.M. Shin, H.W. Jung, J.C. Hyun, J.S. Lee, D. Cho, and Y.L. Joo, J. Appl. Polym. Sci., 120, 752 (2011).
(28.) Z. Zhao, J. Li, X. Yuan, X. Li, Y. Zhang, and J. Sheng, J. Appl. Polym. Sci., 97, 466 (2005).
(29.) M. Bognitzki, W. Czado, T. Frese, A. Schaper, M. Hellwig, M. Steinhart, A. Greiner, and J.H. Wendorff, Adv. Mater., 13, 70 (2001).
(30.) L.H. Catalani, G. Collins, and M. Jaffe, Macromolecules, 40, 1693 (2007).
(31.) H. Tsuji, Macromol. Biosci., 5, 569 (2005).
(32.) J.W. Lu, Z.P. Zhang, X.Z. Ren, Y.Z. Chen, J. Yu, and Z.X. Guo, Macromolecules, 41, 3762 (2008).
(33.) S.C. Wong, A. Baji, and S. Leng, Polymer, 49, 4713 (2008).
Chakrit Thammawong, (1) Sutawan Buchatip, (2) Atitsa Petchsuk, (2) Pramuan Tangboriboonrat, (3) Noppavan Chanunpanich, (4) Mantana Opaprakasit, (5) Paiboon Sreearunothai, (1) Pakorn Opaprakasif
(1) School of Bio-Chemical Engineering and Technology, Sirindhorn International Institute of Technology (SIIT), Thammasat University, Pathum Thani 12121, Thailand
(2) National Metal and Materials Technology Center (MTEC), Thailand Science Park, Pathum Thani 12120, Thailand
(3) Department of Chemistry, Faculty of Science, Mahidol University, Bangkok 10400, Thailand
(4) Industrial Chemistry Department, Faculty of Applied Science, King Mongkut's University of Technology North Bangkok, Bangkok 10800, Thailand
(5) Department of Materials Science, Faculty of Science, Chulalongkorn University, Bangkok 10330, Thailand
|Printer friendly Cite/link Email Feedback|
|Author:||Thammawong, Chakrit; Buchatip, Sutawan; Petchsuk, Atitsa; Tangboriboonrat, Pramuan; Chanunpanich, No|
|Publication:||Polymer Engineering and Science|
|Date:||Feb 1, 2014|
|Previous Article:||Facile preparation of ferroelectric polyvinylidene fluoride-co-trifluoroethylene thick films by solution casting.|
|Next Article:||Computer-Aided Mathematical Modeling and Numerical Simulation and Theoretical Analysis of the Polypropylene Polymer Air Drawing in Spunbonding...|