Effect of Weakly Interacting Nanofiller on the Morphology and Viscoelastic Response of Polyolefins

Transcription

1 Effect of Weakly Interacting Nanofiller on the Morphology and Viscoelastic Response of Polyolefins Jan Kalfus, 1 Josef Jancar, 1 Jaroslav Kucera 2 1 Institute of Materials Chemistry (IMC), School of Chemistry, Brno University of Technology, Brno, Czech Republic 2 Polymer Institute Brno Ltd., Brno, Czech Republic Effect of silica nanofiller on the deformation response and morphology of low- and high-density polyethylene (HDPE, LDPE) and isotactic polypropylene (PP) modified with fumed silica was investigated. The dynamicmechanical thermal spectroscopy, differential scanning calorimetry, optical microscopy, and density measurements were carried out to determine the temperature dependence of storage and loss moduli as well as nanocomposite morphology. It was demonstrated that the degree of matrix reinforcement is considerably affected by the extent of matrix crystallinity, especially, in the temperature range from (T m 1308C) to T m. Based on experimental evidence and literature review, it is proposed that this phenomenon may be attributed to the alpha-mechanical relaxation process occurring above matrix T g. As a result of adding silica into the melted matrix, mobility of chains in contact with silica particles became reduced. This caused substantial changes in morphology of these semicrystalline nanocomposites. POLYM. ENG. SCI., 00: , ª 2008 Society of Plastics Engineers INTRODUCTION So far, serious attempts to investigate the reinforcing mechanism in polymer nanocomposites have been predominantly restricted to systems with amorphous matrices [1 14]. In these matrices above their T g, the chains can undergo segmental immobilization induced even by weak interactions on the large filler-matrix internal interface. It has been clearly demonstrated that the immobilization of polymer chains is the first-order mechanism responsible for the reinforcement in amorphous polymers filled with high specific surface area fillers above the matrix T g. Formation of filler particle network, frequently used as the main source of reinforcement in composites with Correspondence to: Jan Kalfus; kalfus@fch.vutbr.cz Contract grant sponsor: Czech Ministry of Education, Youth and Sports; contract grant number: MSM DOI /pen Published online in Wiley InterScience ( VC 2008 Society of Plastics Engineers submicron-sized fillers, seems to be a second-order effect [4, 13]. In semicrystalline polymer matrices, the high specific surface area filler can cause the immobilization of mobile amorphous phase as well as affect morphology of the crystalline phase. In polymers such as PP and PE, the content of mobile amorphous phase content increases from 0 to 100 wt% nonlinearly in the temperature range from T g to T m. At room temperature, the mobile amorphous phase forms about 5 20 wt% in HDPE or PP [15 18]. Thus, in the case of nanocomposite with polyolefin matrix, the contribution of chain immobilization to the reinforcement should be proportional to the mobile amorphous phase content. This hypothesis seems to be supported by the observed substantially greater increase of the elastic modulus of composites with amorphous matrices compared to the semicrystalline composites with the same volume fraction of high specific surface area filler above the matrix T g [19]. In addition to affecting behavior in the solid state, crystallization kinetics can also be changed by the presence of high specific surface area filler altering the morphology of the crystalline phase [20 23]. This effect can have important consequences in the processing of semicrystalline matrix nanocomposites and their physical ageing. In this contribution, preliminary results are presented showing the effect of adding untreated silica nanoparticles into high-crystallinity HDPE, medium-crystallinity PP, and low-crystallinity LDPE on their viscoelastic response and morphology. EXPERIMENTAL Fumed silica with the specific surface area of 390 m 2 /g and mean particle diameter of 7 nm was used as the filler (Sigma Aldrich, Germany). High-density polyethylene (HDPE) Liten MB 71 (Chemopetrol, Czech Republic) with MFI ¼ 8 g/10 min (1908C, 2.16 kg), low-density polyethylene (LDPE) Bralen NA 7 25 (Slovnaft, Slova- POLYMER ENGINEERING AND SCIENCE -2008

2 kia) with MFI ¼ 7 g/10 min (1908C, 2.16 kg), and isotactic polypropylene (PP) Mosten GB 005 (Chemopetrol, Czech Republic) with MFI ¼ 5 g/10 min (2308C, 2.16 kg) were used as the matrices. To reduce the particle agglomeration, composite samples were prepared using blending in the hot xylene solution. Desired amount of the polymer and filler was added into a beaker with xylene, intensively stirred at 1208C for a few minutes until the whole matrix was dissolved. After the stirring, ultrasonication of composite dispersion was carried out for 5 min employing Sonopuls HD3200 (Bandelin, Germany). The neat polyolefin, used as the reference material, came through the same processing steps as the nanocomposite. Dried nanocomposite samples were compression molded into the mm 3 sheets using laboratory press TP 400 (Fontijne, Netherlands) at 1808C (PE) or 2208C (PP) and clamping force 100 kn for 5 min. Viscoelastic behavior was investigated using the dynamic-mechanical thermal analyzer DX04T (RMI, Czech Republic). Rectangular specimens mm 3 were cut from the compression-molded sheets. The measurements were performed with single cantilever fixture in linear viscoelastic regime under the fixed frequency of f ¼ 5 Hz and a rate of heating of 28C/min within the temperature range from 21208C to the maximum temperature where the specimen remained solid. Differential scanning calorimetry (DSC) measurements were carried out using the DSC Pyris 1 (Perkin Elmer, USA) with a rate of heating and cooling of 108C/min. Before each DSC measurement, heating of the specimen to 1808C in the case of PE or to 2208C in case of PP with 5 min isothermal step at 1808C/2208C was carried out to remove its thermal history. Regular cubes with weight of (5 6 1) mg were used in the DSC measurements. To characterize the effect of silica nanofiller on the density of PE and PP composite, series of pycnometry measurements were carried out at 238C. Presented data are an average values out of three measurements. Standard deviation of less than 5% has been obtained for all the measurements performed. FIG. 2. Storage and loss modulus dependence on temperature for LDPE samples. RESULTS AND DISCUSSION The measured temperature dependences of storage and loss moduli are shown in Figs The three principal contributions responsible for the reinforcement in polymer composites are: (i) substitution of soft polymer matrix by stiff filler, (ii) immobilization of polymer molecules on filler particle surfaces as a consequence of filler polymer interaction, and (iii) stress transfer from the matrix to the filler, providing the filler has nonspherical particles. Moreover, accompanying effects such as particle aggregation, residual stresses, or presence of defects in the polymer matrix or/at the matrix-filler interface can complicate the prediction of elastic moduli. Considering approximately the spherical shape of silica nanoparticles, one may ask which of the two contributions mentioned earlier, (i) substitution or (ii) immobilization, is prevailing in the nanosilica-filled polyolefin within the temperature region T g \ T \ T m. Generally, it can be assumed that the thermal mobility of the backbone chain is negligible regardless the polymer chain is in the crystalline or amorphous phase at temperatures below T g. Hence, adding high specific surface area filler can only FIG. 1. Storage and loss modulus dependence on temperature for HDPE samples. FIG. 3. samples. Storage and loss modulus dependence on temperature for PP 2 POLYMER ENGINEERING AND SCIENCE DOI /pen

3 FIG. 4. Relative storage modulus (related to the neat resin) dependence on temperature for HDPE samples. FIG. 6. Relative storage modulus (related to the neat resin) dependence on temperature for PP samples. affect low temperature secondary transitions and the reinforcement mechanism should primarily be the substitution mechanism (i), where part of soft matrix is replaced by the stiffer filler. In the temperature range T g \ T \ T m, the chains in the amorphous region became mobile as revealed from the NMR or Raman spectra [15 18]. The amount of mobile phase is, in this case, much lower than the amount of the overall amorphous phase suggesting that some chains or their portions form an interphase in vicinity of the surface of crystalline lamellae. Struik [24] has shown that these interphases in semicrystalline polymers and interphases in amorphous polymer composites above the matrix T g exhibit very similar viscoelastic response. The temperature dependence of the relative storage modulus, E 0 r, representing modulus of a composite related to that of the neat matrix is shown in Figs. 4 6 for HDPE, LDPE, and PP samples, respectively. In the case of all three polyolefin matrices, the two regions of the reinforcement described earlier can clearly be distinguished. At low temperature, the E 0 does not exceed values predicted by the simple Kerner-Nielsen micromechanics model [19]. Based on the K-N model, one can predict dependence of the E 0 r on the filler volume fraction, v f, using a set of simple equations as follows [19]: where and E 0 r ¼ 1 þ ABv f 1 Bcv f ; (1) A ¼ð7 5m 1 Þð8 10m 1 Þ 1 (2) B ¼ E 1 f Ef 1 þ A : (3) E m E m The A is the factor depending on the matrix Poisson s ratio, m 1, and B is the factor primarily related to the filler/ matrix stiffness ratio. The factor c represents boundary condition and can be expressed using empirical function taking into account the maximum filler volume fraction, v max, as follows: c ¼ 1 þ v f ð1 v max Þv 2 max : (4) FIG. 5. Relative storage modulus (related to the neat resin) dependence on temperature for LDPE samples. The K-N model represents the substitution mechanism of the composite reinforcement. At temperatures above T g, the E 0 exceeds the K-N micromechanics prediction significantly, being an indication of the gradual onset of an additional reinforcing mechanism the segmental immobilization (see Fig. 7). In the case of the LDPE matrix possessing the lowest degree of crystallinity, the segmental immobilization clearly becomes the primary reinforcing mechanism at temperatures well above the matrix T g. It is in agreement with our assumption mentioned earlier the immobilization DOI /pen POLYMER ENGINEERING AND SCIENCE

4 FIG. 7. Relative storage modulus (related to the neat resin) dependence on temperature for HDPE, LDPE, and PP composite with 8 vol% of silica nanofiller. FIG. 8. Example of determination of a-relaxation contribution to the E 00 temperature dependence. reinforcing mechanism is the most pronounced in lowcrystalinity LDPE and the least pronounced in high-crystalinity HDPE. The medium crystalline PP exhibited behavior close to the HDPE matrix nanocomposite. Additional effects resulting from different supermolecular structure of these polymers have not been considered. Interestingly, the measured relative storage modulus of nanocomposites was below the K-N model-based prediction at temperatures closely below the T g. This phenomenon may have a number of reasons. One possible explanation may be related to the kinetics of the glass transition of the amorphous part of matrix. During the cooling stage of the test as the specimen approaches the T g, polymer chains near the filler particles may be further from the equilibrium response in comparison to the unperturbed bulk chains due to the surface induced segmental immobilization [7]. Described effect may be responsible for the lower density of matrix shell surrounding silica nanoparticle core and causing the reduction of composite modulus in comparison to the K-N based prediction. Another explanation may be based on the assumption that the semicrystalline matrix in the nanocomposite is not of the same structure when compared with the neat polymer resulting in lower modulus of elasticity. This would be possible if only the matrix crystallinity is markedly reduced by interaction of chains with the extent nanoparticle surface. In the temperature region above T g and below T m, the mentioned effect should be overwhelmed by the immobilization phenomenon. In all the E 0 r 2 T dependences, the onset of the E0 deviation from the K-N prediction occurred about C below the melting temperature, T m (see Fig. 7). The actual molecular mechanism responsible for the a- relaxation appearing on the loss modulus temperature dependence of polyolefins above T g has not been fully understood yet [17, 26 31]. Both the amorphous and crystalline portions of polymer matrix are required for the a-relaxation to occur, and it seems highly probable that some kind of constrained reptation of chains through crystallites is essential of this relaxation transition [17]. Nevertheless, regardless the true molecular origin of the a-relaxation one may consider this transition as region, where mobility of polymer segments is largely increased. The loss modulus data for the three polymers investigated are shown in Fig. 8 for the temperature region typical of a-relaxation process. Normalizing the single a-peak in a manner shown in Fig. 9, the onset of the a-relaxation in respect to the melting point was obtained. Surprisingly, the a-peak starts right about 1308C below the melting temperature for all the matrices investigated. It is in accordance with the onset of the contribution of chain immobilization to the overall reinforcement determined from the temperature dependence of the storage modulus, E 0, data (see Fig. 7). In Fig. 10, contributions of the a- relaxation to the loss modulus, a E 00, are depicted. The data suggest that the addition of silica nanofiller induces FIG. 9. Normalized contribution of a-peak to the loss modulus in dependence on normalized temperature for neat HDPE, LDPE, and PP, respectively. Normalization was carried out in respect to value of both maximum of E 00 and melting temperature. 4 POLYMER ENGINEERING AND SCIENCE DOI /pen

5 FIG. 10. Contribution of the a-peak to the loss modulus in dependence on temperature normalized to the peak maximum. some kind of long-range restriction to the overall segmental mobility resulting from the presence of crystalline phase. In this contribution, we will not present detailed analysis of the E 00 T dependences due to the fact that little amount of data is currently available. Deeper understanding of this phenomenon will be a matter of the follow up paper, where the viscoelasticity data will be supplemented with those obtained from the temperature dependent X- ray diffraction and NMR measurements, respectively. During cooling of polyolefin melt below the T m, local volumes having higher degree of molecular order appear in the system as a consequence of fluctuations. These local volumes represent nuclei necessary for the crystallization can occur [31, 32]. During this phase transition, chain mobility is of pivotal importance, since the rate of crystal growth is controlled by the diffusion of segments to the growing surface [33]. In the case of nanocomposite melt, due to the large filler-specific surface area large portion of chains is weakly bounded to the filler surface resulting in lengthening the chain reptation time. In other FIG. 12. Change in the overall degree of crystallinity with v f. words, diffusion of segments to the growing crystal is decelerated. Thus, it seems reasonable to expect the reduced crystallization rate of nanocomposite as compared with the neat matrix [21 23]. In Fig. 11, dependence of the degree of the silica volume fraction is presented. The crystallinity was determined from the enthalpy of the first crystallization peak according to the following equation [20]: X c ¼ DH c 1 wf DH 0 c 1 100; (5) where DH c is the measured enthalpy of crystallization, w f is the filler weight fraction, and DH 0 c is theoretical enthalpy of crystallization for infinitely large crystal equal to 293 J/g for both LD- and HDPE and 207 J/g for PP [34]. One can see that the degree of crystallinity decreased with v f in the case of all three polyolefin matrices. Surprisingly, dependence of the decrease of the degree of crystallinity on the v f was pronounced for LDPE matrix the most (see Fig. 12). Similarly, deviation of composite density from the prediction based on the simple rule of mixtures was observed (see Fig. 13). FIG. 11. Overall degree of crystallinity as a function of v f. FIG. 13. Change in density as a function of v f. DOI /pen POLYMER ENGINEERING AND SCIENCE

6 CONCLUSIONS It was demonstrated that the silica nanofiller was able to significantly affect both the morphology and viscoelastic response of all the polyolefins investigated. In the temperature range T g \ T \ T m, the reinforcing effect of nanoparticles was most pronounced in the LDPE matrix having the lowest degree of crystallinity, supporting the immobilization as a main reinforcement mechanism operating in this temperature region. The onset of the immobilization seems to be connected primarily with the mechanical a-relaxation. Addition of silica nanoparticles also led to the decrease of the overall degree of crystallinity. REFERENCES 1. J. Kalfus and J. Jancar, Polym. Compos., 28, 365 (2007). 2. J. Kalfus and J. Jancar, Polym. Compos., 28, 743 (2007). 3. J. Kalfus and J. Jancar, J. Polym. Sci. Part B: Polym. Phys., 45, 1380 (2007). 4. S.S. Sternstein and A. Zhu, Macromolecules, 35, 7262 (2002). 5. A. Zhu and S.S. Sternstein, Compos. Sci. Technol., 63, 1113 (2003). 6. A. Zhu and S.S. Sternstein, Mater. Res. Soc. Symp. Proc., 661, KK (2001). 7. J. Berriot, H. Montes, F. Lequeux, D. Long, and P. Sotta, Europhys. Lett., 64, 50 (2003). 8. J. Berriot, F. Lequeux, L. Monnerie, H. Montes, D. Long, and P. Sotta, J. Non-cryst. Solids, , 719 (2002). 9. P. Cassagnau, Polymer, 44, 2455 (2003). 10. Q. Zhang and L.A. Archer, Langmuir, 18, (2002). 11. G. Raos, M. Moreno, and S. Elli, Macromolecules, 39, 6744 (2006). 12. S. Kaufman, W.P. Slichter, and D.D. Davis, J. Polym. Sci. Part A-2, 9, 829 (1971). 13. S. Asai, H. Kaneki, M. Sumita, and K.J. Miysaka, J. Appl. Polym. Sci., 43, 1253 (1991). 14. J. O Brien, E. Cashell, G.E. Wardell, and V.J. McBrierty, Macromolecules, 9, 653 (1976). 15. L. Vilc and J. Kratochvila, in Proceedings of the 11th International Conference Polymeric Materials, Halle/Salle, Germany, PI-26 (2004). 16. H.G. Olf and A. Peterlin, Kolloid-Zeitschrift Zeitschrift Polymere, 215, 97 (1967). 17. G. Strobl, The Physics of Polymer, 2nd ed., Springer-Verlag, Berlin (1997). 18. R. Mutter, W. Stille, and G. Strobl, J. Polym. Sci. Part B: Polym. Phys., 31, 99 (1993). 19. L.E. Nielsen and R.F Landel, Mechanical Properties of Polymers and Composite, 2nd ed., Marcel Dekker, New York (1994). 20. Q. Wu, X. Liu, and L.A. Berglund, Macromol. Rapid Commun., 22, 1438 (2001). 21. A. J Waddon and Z. S Petrovic, Polym. J., 34, 876 (2002). 22. K. Nitta, K. Asuka, B. Liu, and M. Terano, Polymer, 47, 6457 (2006). 23. K. Asuka, B. Liu, M. Terano, and K. Nitta, Macromol. Rapid Commum., 27, 910 (2006). 24. L.C.H. Struik, Polymer, 28, 1534 (1987). 25. R.H. Boyd, Macromolecules, 17, 903 (1984). 26. I.M. Ward and J. Sweenly, The Mechanical Properties of Solid Polymers, Wiley, New York (2004). 27. D. Hentschel, H. Sillescu, and H.W. Spiess, Macromolecules, 14, 1605 (1981). 28. K. Schmidt-Rohr and H. W. Spiess, Macromolecules, 24, 5288 (1991). 29. M. L. Mansfield, Macromolecules, 20, 1384 (1987). 30. J.M. Crissman, J Polym Sci Polym Phys Ed., 13, 1407 (1975). 31. P.J. Flory, Proc. R. Soc. London A, 234, 60 (1956). 32. L. Mandelkern, The Physical Properties of Polymers, Cambridge University Press, Cambridge (2003). 33. J.D. Hoffman, Polymer, 23, 656 (1982). 34. G.W. Ehrenstein, G. Riedel, and P. Trawiel, Thermal Analysis of Plastics, Theory and Practice, Hanser, Munich (2004). 6 POLYMER ENGINEERING AND SCIENCE DOI /pen