Fabrication and characterization of biodegradable poly(lactic acid)/layered silicate nanocomposites.
In the past decade significant attention has been focused on the biocompatible and biodegradable polymers, both from ecological and biomedical perspectives . In general, synthetic polymers that are produced from petrochemical products have low recovery/reproduction rates and are not easily degraded in the environment. The rapid growth of municipal waste drives efforts toward biocompatible/biodegradable polymers that can be used as renewable resources for polymer manufacturing and reduce the waste volume of plastics. The most popular and important biodegradable polymers are aliphatic polyesters [2-10], such as poly(lactic acid) (PLA), poly(glycolic acid) (PGA), poly(3-hydroxybutyrate) (PHB), and poly(e-caprolactone) (PCL), among which PLA has attracted the most attention due to its renewable resources , biodegradation, biocompatibility, superior thermal/mechanical properties, and the transparency of the processed materials . Therefore, PLA is widely used in medical application such as wound closure, surgical implants , resorbable sutures , tissue culture , and controlled release systems [7-10].
PLA is a linear aliphatic thermoplastic polyester, produced from renewable resources, and is readily biodegradable through hydrolytic and enzymatic pathways [2, 11-13]. PLAs are synthesized by ring-opening polymerization of lactides or by condensation polymerization of the lactic acid monomers, and those monomers are obtained from the fermentation of corn, potato, sugar cane, and sugar beet . Generally, commercial PLA grades are copolymers of poly(L-lactide) and poly(DL-lactide). The amount of D-enantiomers is known to affect the properties of PLA, such as the melting temperature and the degree of crystallinity. Also, many investigations have been performed to enhance the impact resistance of PLA and to compete with low-cost commodity polymers. These efforts have made use of biodegradable and non-biodegradable plasticizers or by blending PLA with other polymers . From a biomedical point of view the mechanical properties of neat PLA might not be adequate for a high-load-bearing application  which is necessary with additional incorporation of reinforced filler, such as calcium phosphate in the crystalline form of hydroxyapatite (HA)  or clay [18-21].
Polymer layered silicate nanocomposites (PLSNs) have been the focus of academic and industrial attention in recent years because the final composites often exhibit a desired enhancement of physical and/or chemical properties relative to the neat polymer matrix [22-24]. The synthesis of PLSNs is done by the intercalation of monomers or polymers into swellable layered silicate hosts. In most cases the synthesis involves either intercalation of a suitable monomer and then exfoliating the layered host into their nanoscale elements by subsequent polymerization or melt-direct polymer intercalation by using a conventional polymer extrusion process [25-28]. The high aspect ratio layered silicate affects the mechanical, physical, and thermal properties of the synthesizing polymer nanocomposites.
Recently, the enhancement of physical properties of PLA by addition of clay has been extensively reported in the literature [18-21]. The fabrications of PLA/clay nanocomposites were mixed the PLA matrix and organically modified clay using solution intercalation  or melt blending [19-21]. All results indicate that the prepared nanocomposites are mostly intercalated even though with the various surface modifications of clay.
In this report we suggest a new method to prepare the PLA/layered silicate nanocomposites by adding lactic acid (LA) emulsion between the organically modified layered silicate and PLA to enhance the chemical similarity between PLA and layered silicate. First, the layered silicate was modified by n-hexadecyl trimethyl-ammonium bromide (CTAB) cations and then polymerized by a mixture of styrene and methyl methacrylate monomer at a styrene/methyl methacrylate ratio of 8/2 with potassium presulfate as a catalyst. In order to further improve the interaction between the polymer matrix and layered silicate, the LA emulsion was synthesized and mixed with a surface-modified layered silicate. The PLSNs were prepared through the direct insertion of PLA polymer chains from the solution into the surface-treated layered silicate. The morphology and physical properties of prepared nanocomposites were measured by X-ray diffraction (XRD), field-emission scanning electron microscopy (FESEM), polarized optical microscopy (POM), and dynamic mechanical analysis (DMA).
Preparation of PLA/LA/Layered Silicate Nanocomposites
The natural sodium layered silicate with a trioctahedral smectite structure and a cation exchange capacity (CEC) of 110 meq / 100 g was used as the dispersed phase to reinforce the PLA matrix. PLA pellets with a melt index of ~10 g / 10 min were kindly supplied by Wei Mon Industry (Taipei, Taiwan). The surface of natural sodium layered silicate was modified by cationic exchange between [Na.sup.+] in layered silicate galleries and CTAB cations in an aqueous solution at 60[degrees]C for 2 h. The exchanged montmorillonite was then polymerized by a mixture of styrene (ST) and methyl methacrylate (MMA) monomer at a ratio of ST:MMA = 8:2 with potassium presulfate as a catalyst to further improve the interaction between polymer matrix and surface-modified layered silicate. In order to further improve the chemical similarity between the polymer matrix and organic-modified layered silicate, the LA emulsion was prepared and then mixed with surface-modified layered silicate. The PLSNs were prepared using an aqueous surface-treated layered silicate solution mixed with a methylene chloride solution and PLA for a period of time. The PLSNs were then precipitated with an excess amount of deionized water, washed with hot water at 80[degrees]C for 8 h, and dried at 100[degrees]C for 12 h in vacuum.
Powders of pure PLA and 0.1 wt%, 0.3 wt%, 0.5 wt%, 1 wt%, and 2 wt% PLSNs were sandwiched between two cover glasses and heated on a hot stage at 200[degrees]C. The sample was pressed into a thin film with a thickness in the range of 0.10 mm.
FTIR spectroscopy was used to characterize the structure of PLSN. The final spectrum presented is an average of three spectra recorded at different regions over the entire range of the sample. X-ray [theta]/2[theta] diffraction scans of these specimens were obtained using a 3 kW Rigaku III diffractometer equipped with Ni-filtered CuK[alpha] radiation in the reflection mode. Ultrathin sections of the PLSN thin film with a thickness of ~50 nm were prepared with an ultramicrotome equipped with a diamond knife. TEM was carried out with a JEOL transmission electron microscope using an acceleration voltage of 120 keV. Due to the high electron density difference between clay and polymer matrix, staining of the samples was not necessary.
Optical microscopy was carried out using a Zeiss optical microscope equipped with a Mettler FP-82 hot stage and crossed polarizers. The crystallization process was examined using the following temperature sequence. The prepared nanocomposite was heated to melt at 200[degrees]C for 20 min on the hot stage to remove previous thermal history and then cooled quickly to the proposed [T.sub.c]. Optical microscopy was recorded at the proposed [T.sub.c] for various times. The density of nuclei was evaluated by counting the number of spherulite in each photograph per section area. DMA experiments were performed on a Perkin Elmer instrument DMA 7e apparatus equipped with a film tension clamp. The instrument was programmed to measure E' (storage modulus) over a range of 30-150[degrees]C at 2[degrees]C/min heating rate and 1 Hz constant frequency. Calibrations for force, mass, position, and temperature were made in accordance with Perkin Elmer procedures. The specimen films were cut with length-to-width ratios >6 to guarantee that the uniform strain was well within the linear viscoelastic region of the samples and that the collected data were reproducible.
[FIGURE 1 OMITTED]
RESULTS AND DISCUSSION
FTIR spectroscopy was used to characterize the interfacial interaction between LA and organically modified layered silicates. Figure 1 shows the FTIR spectrum of the CTAB-modified layered silicate with and without LA emulsion. The peaks around 1215 and 1746 [cm.sup.-1] corresponding to the C = O and C-O strengthening mode of lactic acid group were obtained . This result indicates that the lactic acid group was successfully grafted to the CTAB-modified layered silicate during the solution mixing process. Because of the chemical similarity between LA emulsion and the PLA polymer chain, CTAB-modified layered silicate with LA emulsion can provide a further interaction between organically modified layered silicates and the PLA matrix.
[FIGURE 2 OMITTED]
[FIGURE 3 OMITTED]
Figure 2 shows the X-ray diffraction data of neat layered silicate and layered silicate modified with a mixture of CTAB, ST, and MMA with and without LA emulsion. It is clear that the X-ray data shifts into a smaller angle for the modified clay. The interlayer distances of the clays were obtained from the peak position ([d.sub.001]-reflection) of WAXD traces. The [d.sub.001]-reflection for the neat layered silicate was found at a 2[theta] [approximately equal to] 7.13[degrees], which corresponds to an interlayer distance of 12.6 [Angstrom] (Fig. 2, trace a). Surface modification by a mixture of surfactant afforded substantially increased interlayer distances of silicate. The X-ray reflection of surface-modified clay (trace b in Fig. 2) was found at 2[theta] of 2.25[degrees], corresponding to an interlayer distance of 39.3 [Angstrom]. Further modification by an LA emulsion slightly increased the interlayer distance of silicate to 40.28 [Angstrom] (Fig. 2, trace c). The X-ray diffraction curves of 0.5 wt%, 1 wt%, and 2 wt% PLSNs are also shown in Fig. 2. All X-ray diffraction data of nanocomposites exhibited no [d.sub.001]-reflection in the relevant region, thus indicating the presence of interlayer distances at least larger than 48 [Angstrom]. These results also show that PLA can be well dispersed in the surface-treated layered silicates, thus further increasing the interlayer distances of silicates. Although X-ray is the simplest method to measure the interlayer distance of silicates, TEM was also used to visually evaluate the degree of intercalation and the amount of aggregation of clay clusters. Figure 3 shows micrographs of TEM of 1 wt% PLSN, in which the gray areas represent the silicate layers in the PLA matrix (bright). From the TEM results the layered silicate is disorderedly intercalated in
the PLA matrix. Therefore, all these results demonstrate that most of the swellable silicate layers are disorderedly intercalated into PLA matrix.
[FIGURE 4 OMITTED]
[FIGURE 5 OMITTED]
Figure 4 shows X-ray diffraction data of PLSNs after melting at 200[degrees]C and then cooled to [T.sub.c] = 100[degrees]C for 1 h. The neat PLA exhibit a very strong crystalline peak at 2[theta] [congruent to] 16.6[degrees]C, corresponding to the (200) and/or (110) plane of a typical orthorhombic crystal. All results of PLSNs show similar X-ray crystalline peaks at 2[theta] [congruent to] 16.6[degrees]C, but the intensities of crystalline peaks decrease with increasing the content of layered silicates. This result indicates that the crystallinity of PLSNs decreases as the content of layered silicate increases. Therefore, we can conclude that the addition of organically modified layered silicate into PLA matrix could not change the crystalline structure of PLA. but the arrangement of PLA polymer chain decreases as the content of layered silicate increases. POM was used to compare the crystal morphologies of neat PLA and PLSNs. Figure 5 presents POM images of PLA and PLSNs crystallized at [T.sub.c] = 100[degrees]C. It is clear that the nucleation density increases and the spherulite texture becomes finer with increasing the content of layered silicates. When PLA was crystallized from the melt, the grown spherulites exhibit typical extinction crosses under the polarizing microscopy. By adding a small amount of layered silicate into PLA matrix, the POM data contains the irregular sheaf-like spherulite and the numbers of spherulites in PLSNs are much higher than that of neat PLA spherulites. For this reason the addition of layered silicate-induced heterogeneous nucleation is attributed to reducing the size of the spherulite.
[FIGURE 6 OMITTED]
[FIGURE 7 OMITTED]
Figure 6 shows the images of SEM of neat PLA matrix and PLSNs. It is clear that the PLA matrix contains an almost uniform morphology of porosity and its diameter is ~3.8 [micro]m. By adding layered silicate into PLA matrix the diameter of porosity of 0.1 wt% PLSN is in the range of 15.3-38.3 [micro]m and increases up to 113-169 [micro]m in the presence of 2 wt% layered silicates in PLSNs. Therefore, all these results demonstrate that the diameter of porosity of PLSNs is much larger than that of neat PLA matrix and increases with increasing the content of layered silicates.
Figure 7 shows the dynamic storage modulus G' of neat PLA matrix and PLSNs over a temperature range of -20 to 100[degrees]C. The data of storage modulus at -20[degrees]C and 60[degrees]C are listed in Table 1. At -20[degrees]C the storage modulus of PLA is 1.28 X [10.sup.9] Pa, which decreases with the increasing temperature; at 60[degrees]C it drops to 0.147 X [10.sup.9] Pa. This is attributed to insufficient thermal energy to overcome the potential barrier for transitional and rotational motions of segments of the polymer molecules in the glassy region, whereas above the glass-transition temperature ([T.sub.g]), the thermal energy becomes comparable to the potential energy barriers to the segmental motions. For 0.1 wt% PLSN, significant enhancement of G' can be seen in the lower temperature range, indicating intercalated 0.1 wt% layered silicate into PLA matrix have a strong influence on the elastic properties of the PLA matrix. Below [T.sub.g], the enhancement of G' is increased ~50% as compared to that of the neat PLA. By adding more layered silicate into PLA matrix, the storage modulus G' drastically decreases with the presence of 1 wt% layered silicate in PLSNs and then continuously decreases with increasing the content of layered silicate to 2 wt%. The result indicates that the reinforcement effect of 0.1 wt% PLSN is predominated by the layered silicate, even though there is a small diameter of porosity. As the addition of more layered silicate into the PLA matrix induces more stiffness in fabricated nanocomposites, the enhancement of the storage modulus is expected to be increased. But as the content of layered silicate increases, the diameter of porosity of PLSN also increases. The reinforcement effect of 1 wt% PLSN is almost cancelled by both factors, as the enhancement of the storage modulus is slightly higher (8%) than that of PLA matrix. But as the addition of more layered silicates causes the increase of porosity of fabricated nanocomposites during the preparation process, the G' of PLSNs decreases as the layered silicate content increases from 1 wt% to 2 wt%. These results clearly indicate that the reinforcement effect of 2 wt% PLSN is predominated by the porosity, even in the presence of layered silicate.
The PLSNs were successfully prepared through the solution insertion of PLA polymer chains from the solution into the surface-modified layered silicates and contained disorderedly intercalated silicate layers within a PLA matrix. Mechanical properties of the 0.1 wt% silicate-containing fabricated nanocomposites observed by dynamic mechanical analysis showed significant improvements in the storage modulus when compared to that of neat PLA matrix. The high content of layered silicate decreases the mechanical properties of fabricated nanocomposites, which is probably due to the presence of a large dimension of porosity in fabricated nanocomposites.
1. Y. Ikada and H. Tsuji, Macromol. Rapid Commun., 21, 117 (2000).
2. H. Tsuji and Y. Ikada, J. Appl. Polym. Sci., 67, 405 (1998).
3. H. Urayama, T. Kanamori, and Y. Kimura, Macromol. Mater. Eng., 287, 116 (2002).
4. M.S. Taylor, A.U. Daniels, K.P. Andriano, and J. Heller, J. Appl. Biomater., 5, 151 (1994).
5. R.A. Jain, Biomaterials, 21, 2475 (2000).
6. A.G. Mikos, M.D. Lyman, L.E. Freed, and R. Langer, Biomaterials, 15, 55 (1994).
7. T.G. Park, S. Cohen, and R. Langer, Macromolecules, 25, 116 (1992).
8. K.R. Kamath and K. Park, Adv. Drug Deliv. Rev., 11, 59 (1993).
9. S.S. Davis, L. Illum, and S. Stolnik, Curr. Opin. Colloid Interface Sci., 1, 660 (1996).
10. U. Edlund and A.C. Albertsson, Adv. Polym. Sci., 157, 67 (2002).
11. T. Iwata and Y. Doi, Macromolecules, 31, 2461 (1998).
12. R. G. Sinclair, J. Macromol. Sci., Pure Appl. Chem., A33, 585 (1996).
13. D. Sawai, K. Takahashi, T. Imamura, K. Nakamura, T. Kanamoto, and S.H. Hyon, J. Polym. Sci. Polym. Phys., 40, 95 (2002).
14. G.B. Kharas, F. Sanchez-Riera, and D.K. Severson, Plastics from Microbes, Hanser, New York (1994).
15. O. Martin and L. Averous, Polymer, 42, 6209 (2001).
16. N.C. Bleach, S.N. Nazhat, K.E. Tanner, M. Kellomaki, and P. Tormala, Biomaterials, 23, 1579 (2002).
17. Y. Shikinami and M. Okuno, Biomaterials, 20, 859 (1999).
18. N. Ogata, G. Jimenez, H. Kawai, and T. Ogihara, J. Polym. Sci. Polym. Phys., 35, 389 (1997).
19. S.S. Ray, P. Maiti, M. Okamoto, K. Yamada, and K. Ueda, Macroimolecules, 35, 3104 (2002).
20. P. Maiti, K. Yamada, M. Okamoto, K. Ueda, and K. Okamoto, Chem. Mater., 14, 4654 (2002).
21. M. Pluta, A. Galeski, M. Alexandra, M.A. Paul, and P. Dubois, J. Appl. Polym. Sci., 86, 1497 (2002).
22. E.P. Giannelis, Adv. Mater., 8, 29 (1996).
23. A. Okada and A. Usuki., Mater. Sci. Eng., C3, 109 (1995).
24. M. Ogawa and K. Kuroda, Bull. Chem. Soc. Jpn., 70, 2593 (1997).
25. G. Lagaly, Appl. Clay Sci., 15, 1 (1999).
26. P.C. LeBaron, Z. Wang, and T.J. Pinnavaia, Appl. Clay Sci., 15, 11 (1999).
27. L. Liu, Z. Qi, and X. Zhu, J. Appl. Polym. Sci., 71, 1133 (1999).
28. N. Ogata, G. Jimenez, H. Kawai, and T. Ogihara, J. Polym. Sci., Polym. Phys., 35, 389 (1997).
Tzong-Ming Wu, Ming-Feng Chiang
Department of Material Science and Engineering, National Chung Hsing University, 250 Kuo Kuang Road, Taichung, Taiwan 402
Correspondence to: T.-M. Wu, e-mail: email@example.com
Contract grant sponsor: National Science Council, Contract grant number: NSC92-2622-E-005-006-CC3.
TABLE 1. Dynamic storage modulus and average diameter of porosity of PLA matrix and PLSNs. Average diameter Storage modulus (MPa) Sample of porosity ([micro]m) (-20[degrees]C) (60[degrees]C) PLA 3.8 1280 147 0.1 wt% PLSN 26.8 1920 (50.0%) 73.2 (-50.2%) 1 wt% PLSN 62 1390 (8.6%) 39.7 (-73.0%) 3 wt% PLSN 141 690 (-46.1%) 34.3 (-76.7%)
|Printer friendly Cite/link Email Feedback|
|Author:||Wu, Tzong-Ming; Chiang, Ming-Feng|
|Publication:||Polymer Engineering and Science|
|Date:||Dec 1, 2005|
|Previous Article:||Weld lines in nylon 6 melt-blended nanocomposites.|
|Next Article:||Hydrocarbon stability of perfluorinated polyether rubbers at elevated temperatures.|