Dimensional stability is critical in many applications. For example, if the layers of a microelectronic chip have different thermal or environmental dimensional stabilities, then residual stresses can develop and cause premature failure. Poor dimensional stability can also cause warping or other changes in shape that affect the function of a material. Nanocomposites provide methods for improving both thermal and environmental dimensional stability. The possible mechanism by which nanofillers can affect the coefficient of thermal expansion (CTE) of a polymer has also been observed in traditional fillers.
The dimension stability of nanocomposites was studied by Zeng and Lee. [58]
Figure 1-17 shows the shape changes of injection molded PS and PS/clay nanocomposites under the aforementioned thermal cycle (50 oC, 1 h; 75 oC, 1 h; 105
oC, 1 h; and 135 oC, 1h). The original sample shape is shown in the first row. Pure PS and the extruded PS/20A (dimethyl dehydrogenated tallow ammonium montmorillonite, 20A) nanocomposite are shown in the second row for comparison.
The third row shows the in-situ polymerized pure PS, PS/20A, and PS/MHABS (2-methacryloyloxyethylhexadecyldimethylammonium bromide, MHABS) nanocomposites. All the nanocomposites contain 5 wt % clay. In the absence of clay, the sample shrank greatly, and the shape became highly irregular. Dimension stability at elevated temperature was improved significantly when 5 wt % of clay was present in the in-situ polymerized nanocomposites, as shown in the third row. The exfoliate PS/MHABS exhibited the best dimensional stability. After the heating cycle, although the sample shrank to a certain extent, the original shape and surface smoothness remained. It is noteworthy that the PS/20A nanocomposite prepared by extrusion
compounding did not show much improvement in dimension stability at elevated temperature, as compared to the in-situ polymerized PS/20A nanocomposite with the same clay content.
Figure 1-17. PS and PS/clay nanocomposites after dimension stability test. Clay loading is 5 wt % for all nanocomposites. [58]
1.6.2 Thermal Stability and Flammability
Delaminated composites have significantly higher degradation temperatures than intercalated nanocomposites or traditional clay composites [61]. Some speculate that this increase in stability is due to the improved barrier properties of the composites. If oxygen cannot penetrate, then it cannot cause oxidation of the resin [62]. In addition, the inorganic phase can act as a radical sink to prevent polymer chains from decomposing. The improved thermal stability of some composites may be limited by the lower thermal stability of alkylammonium ions. For example, in intercalated clay/polystyrene composites, the intercalating agent decomposes at about 250 oC. Bonding the intercalating ion to the polystyrene matrix noticeably improved the thermal stability.
Jin and co-worker investigated thermal property of polymer-clay nanocomposites by TGA and cone calorimetry. [21] The thermal stability of the nanocomposite is enhanced relative to that of virgin polystyrene and this is shown in Figure 1-18. Typically, the onset temperature of the degradation is about 50 oC higher for the nanocomposites than for virgin polystyrene.
Figure 1-18. TGA cures for polystyrene, PS, and the nanocomposites. [21]
One invariably finds that nanocomposites have a much lower peak heat release rate (PHRR) than the virgin polymer. The peak heat release rate for polystyrene and the three nanocomposites are also shown graphically in Figure 1-19. P16-3 means that the nanocompoite was formed using 3 % of P16 clay with polystyrene. The peak heat release rate falls as the amount of clay was increased. The suggested mechanism by which clay nanocomposites function involves the formation of a char that serves as a barrier to both mass and energy transport. [59] It is reasonable that as the fraction of clay increases, the amount of char that can be formed increases and the rate at which heat is released is decreased. There has been a suggestion that an intercalated material is more effective than is an exfoliated material in fire retardancy. [21]
Figure 1-19. Peak heat release rates for polystyrene and the three nanocomposites. [21]
The production of a char barrier must serve to retain some of the polymer and thus both the energy released and the mass loss rate decrease. The amount of smoke evolved, specific extinction area, also decreases with the formation of the
formation of the nanocomposite gives a reduction in smoke, however, the presence of additional clay does not decrease smoke.
1.6.3 Mechanical Properties
The cyclic deformation of PS/MMT nanocomposites as a function of temperature was measured by DMA. The temperature dependence of storage modulus and tanδ were shown in Figure 1-20 and 21, respectively. The storage modulus of PS/MMT nanocomposites were greater than that of pure PS and monotonically increased with the clay content in both the glassy and rubbery regions. However, the improvements in the rubbery region were much greater than those in the glassy region.
This behavior indicates that the restricted segmental motions at the organic-inorganic interface are due to large aspect ratios of the clay platelets, and the polymer chains were also well confined inside the clay galleries at the nanoscale level. [63,64] The storage modulus of PS/MMT-3 was 1.2 times higher than that of pure PS, which is comparable to the earlier reported data (1.4 times improvement). [63] The Tgs of the nanocomposites were estimated from the peak values of tanδ in Figure 1-18, which were shifted towards higher temperature with increasing the clay content. These results indicate that nanoscale clay platelets strongly restrict the polymer segmental motions, resulting in the significant increase in Tg. This improvement in Tg is higher than those of other researchers even though the smaller clay content was used in this experiment. [65,66]
Figure 1-20. Storage modulus of (a) pure PS, (b) PS/MMT-1, (c) PS/MMT-2 and (d) PS/MMT-3.
Figure 1-21. Tanδ values of (a) pure PS, (b) PS/MMT-1, (c) PS/MMT-2 and (d) PS/MMT-3.
The effects of clay loadings on tensile properties of the PS/MMT nanocomposites are shown in Figure 1-22. The tensile strength and Young’s modulus were significantly enhanced in the presence of the small contents of clay, while the elongation at break was reduced with increasing the clay content. The increase in tensile strength was attributed to the stronger interfacial adhesion between PS and the clay platelets. However, the enhancement of modulus was reasonably ascribed to the high resistance exerted by the clay platelets against the plastic deformation and the stretching resistance of the oriented polymer backbones in the galleries. The improvement of tensile strength in PS/MMT-3 compared to pure PS was ~47%, which is greater than the earlier reported value in the literature (~21%) for PS/MMT nanocomposite with 3wt% MMT prepared by melt blending. [64] Similarly, the enhancement of Young’s modulus in PS/MMT-3 compared to pure PS was ~25%, which is much greater than the reported value (7.4% improvement for PS/MMT nanocomposite with 5wt% clay prepared by emulsion polymerization). [67] However, the elongations at break were reduced with increasing the clay content. Similar results were earlier reported. For example, the reduction of elongation at break in PS/MMT nanocomposite with 4.4wt% MMT prepared by melt blending was reported to ~26%.
[63]
Figure 1-22. (a) Tensile strengths, (b) Young’s modulus and (c) elongations at break of PS/MMT nanocomposites. [68]
1.6.4 Gas barrier properties
Clays are believed to increase the barrier properties by creating a maze or
“tortuous path” (Figure 1-23) that retards the progress of the gas molecules through the matrix resin. The direct benefit of the formation of such a path is clearly observed in polyimide/clay nanocomposites by dramatically improved barrier properties, with a simultaneous decrease in the thermal expansion coefficient. [78,79] The polyimide/layered silicate nanocomposites with a small fraction of OMLS exhibited reduction in the permeability of small gases, e.g. O2, H2O, He, CO2, and ethylacetate vapors. [80] For example, at 2 wt % clay loading, the permeability coefficient of water vapor was decreased ten-fold with synthetic mica relative to pristine polyimide.
By comparing nanocomposites made with layered silicates of various aspect ratios, the permeability was seen to decrease with increasing aspect ratio.
Figure 1-23 Formation of tortuous path in PLS nanocomposites.
Apparently, significant improvements in barrier properties are also achievable with nonplate-like nanoparticles. [81] Nano Material Inc. reports that a PVA/EVOH matrix composite with 7 nm silica and titania nanoparticles exhibits a gas permeability of 1 cc m-2 d-1 atm-1 and moisture permeability of less than 1 g m-2 d-1. Although this is achieved at very high loadings, the material is melt processable.
The absorption of water into composites is significant. For example, one of the limitations of Nylon is the reduction in mechanical properties that accompanies the
absorption of water. The addition of exfoliated montmorillonite increases the resistance to water permeation after 30 min from 2% to 1% at 5 wt. % of filler. [82]
The mechanism of the reduction is attributed to the constrained region of the Nylon. If the constrained region is taken into account, the diffusion coefficient follows a rule of mixtures. Figure 1-24 shows the change in diffusion coefficient of water in Nylon in response to clay content.
Figure 1-24 Dependence of diffusion coefficient of water on clay content for montmorillonite with a layer width of 100 nm, and saponite with a layer width of 50
nm. [82]
1.6.5 Electrical and Optical Properties
The electrical and optical properties of nanofilled polymers are exciting areas of research. This is particularly true because of the possibility of creating composites with unique combinations of functionalities, such as electrically conducting composites with good wear properties that are optically clear. Such properties can result because nanoparticles, with diameters distinctly below the Rayleigh scattering limit, still display their solid-state physical properties when embedded in transparent matrices.
Optical composites have been defined as composites consisting of optically
active nanoparticles embedded in a transparent host material, often a polymer. Optical composites take advantage of the optical properties of materials that are hard to grow in single-crystal form or that require protection from the environment and give them the ease of processing afforded many polymers. In addition, sometimes the material must be used at the nanoscale to achieve specific optical properties, and the matrix is used just to hold the particles together and provide processability. For example, high-grade optical composites, with properties otherwise obtainable only in optical glasses, become accessible through the use of polymer molding techniques.
1.7 Summary
The nanocomposite presented here is a composite material reinforced with silicate sheets. Silicate sheet is an ultrafine filler of nanometer size, which is almost equal to the size of the matrix polymer. Although the content of the filler is as little as several wt%, individual filler particles exist at a distance as close as tens of nanometers from each other because of their ultrafine size. One end of the polymer is strongly restrained to the silicate sheet by polar interaction. Thus, the nanocomposite has a microstructure that has never been seen in conventional composites. The characteristic properties of the nanocomposite are derived from this very structure.
Considering the properties, the nanocomposite may be, in a sense, an embodiment of the ideal polymer composite, or a completely novel composite. Silicate sheet can be regarded as a rigid inorganic polymer. In this sense, the nanocomposite realized is a molecular composite in which a silicate sheet is used instead of an organic rod-like polymer.
References
[1] Blumstein, A. Bull. Chim. Soc. 1961, 899.
[2] Usuki, A.; Kojima, Y.; Kawasumi, M.; Okada, A.; Fukushima, Y.; Kurauchi, T.;
Kamigaito, O. J. Mater. Res. 1993, 8, 1179.
[3] Lan, T.; Pinnavaia, T. J. Chem. Mater. 1994, 6, 2216.
[4] Usuki, A.; Kato, M.; Okada, A.; Kurauchi, T. J. App. Polym. Sci. 1997, 63, 137.
[5] Jeon, H. G.; Jung, H. T.; Lee, S. D.; Hudson, S. Polymer Bulletin 1998, 41, 107.
[6] Giannelis, E. Adv. Mater. 1996, 8, 29.
[7] Fisher, H.; Gielgens. L.; Koster, T. Nanocomposites from Polymers and Layered Minerals; TNO-TPD Report, 1998.
[8] Carrado, K. A.; Langui, X. Microporous Mesoporous Mater. 1999, 27, 87.
[9] Friedlander, H. Z.; Grink, C. R. J. Polym. Sci., Polym. Lett. 1964, 2, 475.
[10] Blumstein, A. J. Polym. Sci., Part A. 1965, 3, 2653.
[11] Kato, C.; Kuroda, K.; Takahara, H. Clay and Clay Minerals 1981, 29, 294.
[12] Kelly, P.; Moet, A.; Qutubuddin, S. J. Mater. Sci. 1994, 29, 2274.
[13] Akelah, A.; Kelly, P.; Qutubuddin, S.; Moet, A. Clay Minerals 1994, 29, 169.
[14] Akelah, A.; Moet, A. J. Mater. Sci. 1996, 31, 3189.
[15] Akelah, A.; Moet, A. Mater. Lett. 1993, 18, 97.
[16] Giannelis, E. P.; Adv. Mater. 1996, 8, 29.
[17] Vaia, R.; Ishii, H.; Giannelis, E. Chem. Mater. 1993, 5, 1694.
[18] Doh, J. G.; Cho, I. Polymer Bulletin 1998, 41, 511.
[19] Takekoshi, T.; Fouad, F.; Mercx, F. P. M.; De Moor, J. J. M. US patent 5,773,502.
Issued to General Electric Co., 1998.
[20] Okada, K. (Sekisui) Japan Patent 11-228748, 1999.
[21] Zhu, J.; Alexander, B. M.; Frank, J. L.; Charles, A. W. Chem. Mater. 2001, 13, 3774
[22] Dongyan, W.; Jin, Z.; Qiang, Y.; Charles, A. W. Chem. Mater. 2002, 14, 3837.
[23] Dongyan, W.; Charles, A. W. Polym. Deg. Stab. 2003, 82, 309.
[24] Fu, X.; Qutubuddin, S. Polymer 2001, 42, 807.
[25] Lin, J. J.; Cheng, I. J.; Chou, C. C. Macromol. Rapid Commun. 2003, 24, 492.
[26] Yei, D. R.; Kuo, S. W.; Fu, H. K.; Chang, F. C. Polymer 2005, 46, 741.
[27] Yei, D. R.; Kuo, S. W.; Su, Y. C.; Chang, F. C. Polymer 2004, 45, 2633.
[28] Tseng, C.R.; Wu, J.Y.; Lee, H.Y.; Chang, F.C. Polymer 2001; 42: 10063.
[29] Tseng, C.R.; Lee, H.Y.; Chang, F.C. J. Polym Sci Part B: Polym Phys 2001; 39:
[35] Fan, X.; Zhou, Q.; Xia, C.; Crsitopholi, W.; Mays, J.; Advincula, R. C. Langmuir 2002, 18, 4511.
[36] Fan, X.; Xia, C.; Advincula, R. C. Langmuir 2003, 19, 4381.
[37] Meier, L.; Shelden, R.; Caseri, W.; Suter, U. Macromolecules 1994, 27, 1637.
[38] Pruker, O.; Ruhe, J. Macromolecules 1998, 31, 592.
[39] Pruker, O.; Ruhe, J. Macromolecules 1998, 31, 602.
[40] Yano, K.; Usuki, A.; Okada, A.; Kuraychi, T.; Kamigaito, O. Polym. Prepr. 1991, 32, 65.
[41] Kojima, Y.; Usuki, A.; Kawasumi, M.; Okada, A.; Kurauchi, T.; Kamigaito, O. J.
[42] Liu, L.; Qi, A.; Zhu, X. J. Appl. Poly. Sci. 1999, 71, 1133.
[43] Yano, K.; Usuki, A.; Okada, A.; Kuraychi, T.; Kamigaito, O. J. Polym. Sci., Part A: PolymChem. 1993, 31, 2493.
[44] Yano, K.; Usuki, A.; Okada, A. J. Polym. Sci., Part A: Polym. Chem. 1997, 35, 2289.
[45] Lan, T.; Kaviratna, P. D.; Pinnavaia, T. J. Chem. Mater. 1994, 6, 573.
[46] Tyan, H.L.; Liu, Y.C.; Wei, K.H. Polymer 1999, 40, 4877.
[47] Kawasumi, M.; Hasegawa, N.; Kato, M.; Usuki, A.; Okada, A. Macromolecules 1997, 30, 6333.
[48] Kato, M.; Usuki, A.; Okada, A. J. Appl. Polym. Sci. 1997, 66, 1781.
[49] Kuchta, F.D.; Lemstra, P. J.; Kellar, A.; Batenburg, L. F.; Fischer, H. R. MRS Symp. Proc. 1999, 576, 363.
[50] Lee, D. C.; Lee, J. W. J. Appl. Polym. Sci. 1996, 61, 1117.
[51] Okamoto, M.; Morita, S.; Taguchi, H.; Kim, Y.H.; Kotaka, T.; Tateyama, H.
Polymer 2000, 41, 3887.
[52] Chen, G.; Chen, X.; Lin, S.; Ye, W. J. Mater. Sci. Lett. 1999, 18, 1761.
[53] Dietsche, F.; Thomann, Y.; Thomann, R.; Mulhaupt, R. J. Appl. Polym. Sci. 2000, 75, 396.
[54] Hasegawa, N.; Okamoto, H.; Kawasumi, M.; Usuki, A. J. Appl. Poly. Sci. 1999, 74, 3359.
[55] Lan, T.; Pinnavaia, T. J. Chem. Mater. 1994, 6, 2216.
[56] Wang, Z.; Lan, T.; Pinnavaia, T. J. Chem. Mater. 1996, 8, 2200.
[57] Massam, J.; Pinnavaia, T. J. MRS Symp. Proc. 1998, 20, 223.
[58] C. Zeng, L. J. Lee, Macromolecules 2001, 34, 4098.
[59] Gilman, J. W.; Jackson, C. L.; Morgan, A. B.; Harris, R.; Giannelis, E. P.;
Phillips, S. H. Chem. Mater. 2000, 12, 1866.
[60] T. Agag, T. Koga, T. Takeichi, Polymer 2001, 42, 3399.
[61] J. Lee, E. Giannelis, Polym. Prepr. 1997, 38, 688.
[62] S. D. Burnside, E. P. Giannelis, Chem. Mater. 1995, 7, 1597.
[63] Hasegawa, N.; Okamoto, H.; Kawasumi, M.; Usuki, A. J. Appl. Polym. Sci. 1999, 74, 3359.
[64] Park, C.I.; Choi, W.M.; Kim, M.K.; Park, O.O. J. Polym. Sci. Part B: Polym.
Phys. 2004, 42, 1685.
[65] Fu, X.; Qutubuddin, S. Mater. Lett. 2000, 42, 12.
[66] Kim, Y.K.; Choi, Y.S.; Wang, K.H.; Chung, I. J. Chem. Mater. 2002, 14, 4990.
[67] Noh, M.W.; Lee, D.C. Polym. Bull. 1999, 42, 619.
[68] Uthirakumar, P.; Song, M.K.; Nah, C.; Lee, Y.S. Euro. Polym. J. 2005, 41, 211.
[69] Aranda, P.; Ruiz-Hitzky, E. App. Clay Sci. 1999, 15, 119.
[70] Chen, W.; Xu, Q.; Yuan, Z. J. Mater. Sci. Lett. 1999, 18, 711.
[71] Vaia, R. A.; Vasudevan, S.; Krawwiec, W.; Scanlon, L. G.; Giannelis, E. P. Adv.
Mater. 1995, 7, 154.
[72] Kojima, Y.; Fukumori, K.; Usuki, A.; Kurauchi, T.; J. Mater. Sci. Lett. 1993, 12, 889.
[73] Kresge, E. N.; Lohse, D. J. US Patent 1996, 5, 576, 372.
[74] Laus, M.; Francescangeli, O.; Sandrolini, F. J. Mater. Res. 1997, 12, 3134.
[75] Oriakhi, C. O.; Zhang, X.; Lerner, M. M. Appl. Clay Sci. 1999, 15, 109.
[76] Huang, X.; Lewis, S.; Brittain, W. J.; Vaia, R. A. Macromolecules 2000, 33, 2000.
[77] Wu, Q.; Xue, Z.; Qi, Z.; Wang, F. Polymer 2000, 41, 2029.
[78] Yano, K.; Usuki, A.; Okada, A.; Kurauchi, T.; Kamigaito, O. Synthesis and properties of polyimide–clay hybrid. Polym Prepr (Jpn) 1991, 32(1), 65.
hybrid films. J Polym Sci, Part A: Polym Chem 1997, 35, 2289.
[80] Yano, K.; Usuki, A.; Okada, A.; Kurauchi, T.; Kamigaito, O. Synthesis and properties of polyimide–clay hybrid. J Polym Sci, Part A: Polym Chem 1993, 31, 2493.
[81] Modern Plastics, February, 1999.
[82] Kojima, Y.; Usuki, A.; Kawasumi, M.; Okada, A.; Kurauchi, T.; Kamigaito, O. J.
Appl. Polym. Sci. 1993, 49, 1259.
[83] US Patent 4,739007 (19 Apr 1988).
[84] Lee, D. C.; Jan, L. W. J. Appl. Polym. Sci. 1996, 61, 1117.