Apr 10, 2009 - 3.1 Introduction. Over the last few years, the utility of inorganic nanoscale particles as filler to enhance the polymer performance has been ...
3 Rheology in Polymer/Clay Nanocomposites: Mesoscale Structure Development and Soft Glassy Dynamics Masami Okamoto Contents 3.1 Introduction................................................................................................... 57 3.2 Linear Viscoelastic Properties..................................................................... 58 3.3 Nonlinear Shear Response..........................................................................64 3.4 Analogy to Soft Colloids..............................................................................65 3.5 Reversibility of Network Formation Process............................................ 66 3.6 Alignment of Silicate Layers in Networks................................................ 70 3.7 Summary........................................................................................................ 75 References................................................................................................................ 76
3.1 Introduction Over the last few years, the utility of inorganic nanoscale particles as filler to enhance the polymer performance has been established. Of particular interest is recently developed nanocomposite technology consisting of a polymer and organically modified layered filler (organo-clay), because they often exhibit remarkably improved mechanical and various other materials properties as compared with those of virgin polymer or conventional composite (micro-/macrocomposites) [1–6]. These concurrent property improvements are well beyond what can generally be achieved through the micro-/macrocomposites preparation. The synthetic strategy and molecular design was first explored by Toyota group with nylon 6 as the matrix polymer [7]. This new class of material is now being introduced in structural applications, such as gas barrier film and other load-bearing applications [5]. Polymer/clay nanocomposites (PCNs) and their self-assembly behaviors have recently been approached to produce nanoscale polymeric materials [1–6]. Additionally, these nanocomposites
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have been proposed as model systems to examine polymer structure and dynamics in confined environments [8–10]. In order to understand the processability of these materials (i.e., the final stage of any polymeric material), one must understand the detailed rheological behavior of these materials in the molten state. Understanding the rheological properties of PCN melts is not only important in gaining a fundamental knowledge of the processability but is also helpful in understanding the structure–property relationships in these materials. Although rheological measurement is an indirect probe, it is a well-established approach to probe the interaction between nanofiller and polymer matrix and the time-dependent structure development. In addition, more clear nanoscale and mesoscale structure development of the systems could be provided when combined with X-ray/light-scattering experiments and electron microscopy. The original mesoscale structure in PCNs consists of randomly oriented exfoliated layers or tactoids of layers. This randomly distributed nanofiller forms a “organo-clay network” structure that is mediated by polymer chains and organo-clay–organo-clay interactions, responsible for the linear viscoelastic response observed in PCN melts. This mesostructure, which is intrinsically in a disordered metastable state and out of equilibrium, and offers an apt analogy to soft colloidal glasses and gels, was extensively discussed [11,13–27]. In this chapter, we survey mesostructure development in PCN melts with a primary focus on flow behavior and an analogy to soft glassy dynamics [12].
3.2 Linear Viscoelastic Properties Dynamic oscillatory shear measurements of polymeric materials are generally performed by applying a time-dependent strain of γ (t ) = γ o sin(ωt ) , and the resultant shear stress is σ(t ) = γ o [G′ sin(ωt ) + G′′ cos(ωt )] , with G′ and G′′ being the storage and loss modulus, respectively. Generally, the rheology of polymer melts strongly depends on the temperature at which measurement is carried out. It is well known that for thermorheological simplicity, isotherms of storage modulus (G′(ω ) ), loss modulus (G′′(ω )), and complex viscosity (|η*|(ω)) can be superimposed by horizontal shifts along the frequency axis: bTG′( aTω , Tref ) = bTG′ (ω , T ) bTG′′( aTω , Tref ) = bTG′′ (ω , T )
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|η*|( aTω , Tref ) =|η*| (ω , T )
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where aT and bT are the frequency and vertical shift factors, respectively; and Tref is the reference temperature. All isotherms measured for pure polymer and for various PCNs can be superimposed along the frequency axis. In the case of polymer samples, it is expected, at the temperatures and frequencies at which the rheological measurements were carried out, that the polymer chains should be fully relaxed and exhibit characteristic homopolymer-like terminal flow behavior (i.e., curves can be expressed by a power law of G′(ω) ∝ ω2 and G″(ω) ∝ ω). The rheological properties of in situ polymerized nanocomposites with end-tethered polymer chains were first described by Krisnamoorti and Giannelis [13]. The flow behavior of poly(_-caprolactone) (PCL)- and nylon 6-based nanocomposites differed extremely from that of the corresponding neat matrices, whereas the thermorheological properties of the nanocomposites were entirely determined by that behavior of matrices [13]. The slope of G′(ω) and G″(ω) versus the aTω is much smaller than 2 and 1, respectively. Values of 2 and 1 are expected for linear monodispersed polymer melts and are a large deviation, and especially in the presence of a very small amount of layered silicate, loading may be due to the formation of network structure in the molten state. However, such nanocomposites based on the in situ polymerization technique exhibit fairly broad molar mass distribution of the polymer matrix, which hides the structurally relevant information and impedes interpretation of the results. To date, the melt state linear dynamic oscillatory shear properties of various kinds of PCNs have been examined for a wide range of polymer matrices, including nylon 6 with various matrix molecular weights [14], PS [15], polystyrene (PS)-polyisoprene (PI) block copolymers [16,17], PCL [18], polypropylene (PP) [11,19–22], polylactide (PLA) [23,24], and poly(butylene succinate) (PBS) [25,26]. In the linear viscoelastic regime, a big change in the terminal (low-frequency) region from a liquid-like response to a solidlike response was observed for all PCNs (G′(ω) ~ G″(ω) ∝ ω 0), ascribed to the formation of a volume spanning mesoscale organo-clay network above the mechanical percolation threshold [22]. The terminal rheology is sensitive to organo-clay loading and the extent of exfoliation/intercalation in the polymer matrix. A typical example of the linear dynamic viscoelastic master curves for the neat polylactide (PLA) and various PLA-based nanocomposites (PLACNs) with different organo-clay loading [24] is shown in Figure 3.1. The linear dynamic viscoelastic master curves were generated by applying the timetemperature superposition principle and shifted to a common temperature Tref using both frequency shift factor aT and modulus shift factor bT. The moduli of the PLACNs increase with increasing clay loading at all frequencies ω. At high ωs, the qualitative behavior of G′(ω) and G″(ω) is essentially the same and is unaffected by frequencies. However, at low frequencies, G′(ω) and G′(ω) increase monotonically with increasing organo-clay content. In the low-frequency region, the curves can be expressed by power law of G′(ω) ∝ ω2
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105 0.5
bTG´/Pa
104 103
1
102
2
101
PLA PLACN4 PLACN5 PLACN7
100 10–1 Tref = 175°C 104
bTG˝/Pa
1 2
103
102
|η*|/Pa.s
101
104
103 10–2
10–1
100
101
102
ω/rad.s–1
aT
Figure 3.1 Reduced frequency dependence of storage modulus, loss modulus, and complex viscosity of neat polylactide (PLA) and various PLA-based nanocomposites (PLACNs) under dynamic oscillatory shear measurement. (Reprinted from Sinha Ray S, Yamada K, Okamoto M, Ueda K, Polymer 44, 6631. Copyright 2003 Elsevier Science. With permission.)
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Table 3.1 Terminal Slopes of G′ and G″ versus aTω for Polylactide (PLA) and Various PLA-Based Nanocomposites (PLACNs) System
Organo-Clay/wt%
G′
G″
Ea a/kJmol–1
PLA PLACN4 PLACN5 PLACN7
0 4.0 5.0 7.0
1.6 0.2 0.18 0.17
0.9 0.5 0.4 0.32
170 ± 5 225 ± 10 225 ± 5 230 ± 8
Flow activation energy obtained from an Arrhenius fit of master curves. Source: Reprinted from Sinha Ray S, Yamada K, Okamoto M, Ueda K, Polymer 44, 6631, © 2003, Elsevier Science. (With permission.) a
and G″(ω) ∝ ω for neat PLA, suggesting this is similar to those of the narrow molecular weight (Mw) distribution homopolymer melts. On the other hand, for aTω < 5 rad.s–1, the viscoelastic response (particularly G′(ω)) for all the nanocomposites displays significantly diminished frequency dependence as compared with the matrices. In fact, for all PLACNs, G′(ω) becomes nearly independent at low aTω and exceeds G″(ω), characteristic of materials exhibiting a pseudo-solid-like behavior. The terminal zone slopes values of both neat PLA and PLACNs are estimated at the lower aTω region (
