Diblock Copolymer Based Composites for Tire Tread Applications with

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Letter

Diblock Copolymer Based Composites for Tire Tread Applications with Improved Filler Network Topology Cecilia Aguiar da Silva, Sriharish M. Nagaraja, Marc Weydert, and Mario Beiner ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.7b00313 • Publication Date (Web): 16 Feb 2018 Downloaded from http://pubs.acs.org on February 18, 2018

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ACS Applied Nano Materials

Diblock Copolymer Based Composites for Tire Tread Applications with Improved Filler Network Topology C. Aguiar da Silva,† S. M. Nagaraja,‡ M. Weydert,† and M. Beiner∗,‡ Goodyear Innovation Center Luxembourg, L-7750 Colmar-Berg, Luxembourg, and Fraunhofer Institut für Mikrostruktur von Werkstoffen und Systemen IMWS, Walter-Hülse-Str. 1, 06120 Halle (Saale), Germany E-mail: [email protected]

Abstract We present in this letter the results of a study focusing on vulcanized polybutadienestyrene butadiene rubber (PB-SBR) diblock copolymers in the microphase-separated state filled with various amounts of silica nanoparticles. It is demonstrated that the microphase-separation of PB and SBR blocks is preserved in composites with a silica mass content ≤60 parts per hundred rubber (phr). Only effect of silica incorporation is a more restricted long range order of the block copolymer morphology. Most interesting finding is, however, that through mechanical mixing the silica nanoparticles are highly selectively incorporated in the SBR phase up to high filler contents. This opens new opportunities towards a rational design of the filler network in composites with self-assembled elastomer matrices and results in advantageous mechanical properties ∗

To whom correspondence should be addressed Goodyear Innovation Center Luxembourg, L-7750 Colmar-Berg, Luxembourg ‡ Fraunhofer Institut für Mikrostruktur von Werkstoffen und Systemen IMWS, Walter-Hülse-Str. 1, 06120 Halle (Saale), Germany †

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for tire tread applications.

Keywords: rubber composites, block copolymers, filler network, relaxation dynamics, tire treads, dissipation, reinforcement

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Block copolymer composites are considered as perspective for tailored functional materials in various fields of application. 1–5 Interestingly, a similar approach can be also used to optimize the mechanical performance of rubber composites for tire treads. Recently, we have demonstrated based on different series of vulcanized PB-SBR diblock copolymers how relaxation dynamics and dissipation behavior of dual-phase self-assembled elastomers can be influenced by varying their molecular architecture. 6 Styrene content and vinyl fraction will determine the glass transition temperature of both blocks while phase morphology and fraction of interfacial material are of major importance for the shape of the relaxation spectrum. By synthesizing PB-SBR diblock copolymers with appropriate fraction of each block, microstructure (vinyl and cis/trans content) of butadiene sequences and styrene content in the SBR block the application-relevant dissipation behavior of these elastomers can be nicely fine-tuned. A special advantage of block copolymers is here that the phase morphology on the nanoscale is defined by the molecular architecture and only to a minor extent influenced by the processing conditions. This is a real advantage compared to commonly used elastomer blends where the processing conditions are of major importance for phase morphology and resulting performance. 7 The independence of the morphology on the mixing process observed for strongly segregated block copolymers allows to control the frequency temperature position where dissipation of the elastomer matrix occurs. 6,8 In particular, the dissipation in the wet grip relevant range - where dissipation in the tire treads is required during braking - can be optimized well based on the properties of the elastomer matrix. 9,10 A lab indicator which is often used as a first approximation of wet grip is the loss tangent from dynamic mechanical measurements at 0 ◦ C and 10 Hz. The loss tangent in this range should be high for rubber composites used in tire treads. Choosing a diblock copolymer matrix with the right composition of the SBR block this can be easily achieved. However, there is a second range where the frequency-temperature dissipation is extremely important in case of tire treads. A commonly used lab indicator for the rolling resistance is the loss tangent at 60 ◦ C and 10 Hz. In this range the loss tangent should be as low as possible

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in order to save fuel. In elastomer nanoparticle composites this parameter is dramatically increasing with increasing filler content. Based on recent findings one can conclude that this effect is mainly due to dissipation in glassy rubber bridges connecting neighbored filler particles. 11 It has been shown for various types of rubber composites that these glassy rubber bridges are part of the so-called filler network - being a percolating solid structure in highly filled elastomer nanoparticle composites - and responsible for its visco-elastic properties 11–14 (although competing results have been reported in other cases 15 ). Small deformations of these glassy bridges seem to cause a significant part of the dissipation contributions in the rolling resistance relevant range although the total volume fraction of the immobilized rubber is only a very few percent. 16 From a tire application point of view, the filler network is urgently required in order to ensure reasonable values of hardness and mechanical modulus in case of elastomer nanoparticle composites and therefore appropriate tire tread abrasion. 17 Hence, it seems to be an interesting goal to reach high reinforcement with as small as possible filler contents. This task is connected with the aim to minimize the percolation threshold φC which is the minimum filler content where a percolating filler network is formed. One possible approach in lowering the percolation threshold is to fine-tune the filler network topology. Accordingly, a central aim of the presented study was to understand to what extent microphase-separated block copolymers can really be used to influence the filler network topology in elastomer nanoparticle composites. Results from AFM demonstrating that silica nanoparticles are selectively incorporated in the SBR phase of these block copolymers will be presented. Further, it will also be shown that for a given filler content higher reinforcement can be achieved for block copolymers with lamellar morphology as compared to similar blend based composites where high filler fractions are more uniformly distributed in the matrix for application relevant filler fractions (≥ 60 phr). Moreover, it will be shown that the filler localization in the SBR phase of block copolymers is accompanied by a reduced dissipation in frequency temperature range which is relevant for the rolling resistance of tire treads. Atomic force microscopy (AFM) images of a vulcanized PB-SBR diblock copolymer sam-

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Figure 1: AFM phase images for (a) a vulcanized PB-SBR diblock copolymer with lamellar morphology and composites where this dual-phase, self-assembled elastomer is filled with (b) 20 phr, (c) 40 phr and (d) 60 phr silica nanoparticles. The PB domains appear in brown while orange color is indicating SBR domains. The white spots in the composites are filler particles and aggregates. A certain yellow region for the unfilled samples in part (a) is due to additives diffusing to the surface after sample cryo-microtoming.

ple and composites where this self-assembled two-phasic elastomer is filled with different amounts of silica nanofiller (Zeosil 1165MP) are presented in Figure 1. It is nicely seen in Figure 1a that this diblock copolymer shows a lamellar morphology. This is an expected finding considering that the chosen PB-SBR diblock copolymer is nearly symmetric (SBR volume fraction φSBR = 0.46) and has in the relevant temperature range an interaction parameter for which microphase separation should occur (temperature-dependent values for the product of Flory-Huggins interaction parameter χ and degree of polymerization N are χN25◦ C = 61.3 > χN150◦ C = 47.2 > (χN )C = 10.5). 6 The layer spacing is about 80 nm as obtained from small angle x-ray scattering experiments on non-crosslinked samples. 6 Interestingly, there are no indications for a significant change of the block copolymer morphology after vulcanization with a commonly used vulcanization system at 150 ◦ C as seen in Figure 1a. Further details about the block copolymer microstructure and composite processing are given in the Supporting Information. Even more interesting is that this lamellar morphology is basically also preserved if PB-SBR block copolymers are filled with different silica nanoparticle loadings (Figure 1 b-d). A lamellar-like morphology of the block copolymer matrix seems to be still present although the long range order is more and more disturbed with increasing filler content. What obviously happens is that individual lamellae are more 5



often disrupted. Average periodicities taken from a Fast Fourier Transformation (FFT) of AFM images vary only by 5% for different filler contents and range from 80 nm to 77 nm for 0 to 40 phr silica loading, respectively. Note that the thickness of the individual SBR lamellae (about 37 nm are calculated based on periodicity and φSBR , 27 to 30 nm are interactively determined from AFM images for the ’pure’ SBR domains in regions without filler, cf. SI) is maximum twice larger than the average diameter of the primary silica particles Zeosil 1165MP of about 20 nm. 18 Most interesting finding in the AFM images of silica-filled block copolymers is, however, that the silica nanoparticles are preferentially located in the SBR domains up to filler contents of at least 60phr. A more detailed analysis of the AFM images shows that practically all silica particles as surrounded by SBR. This might be somehow expected since there should be a stronger attractive interaction between SBR and silica as compared to PB due to the polarity of the blocks. However, the high degree of selectivity seems to be surprising considering recent results for elastomer blends where selectivity effects are seen for small filler contents 19–21 but usually much less pronounced for high filler loadings. AFM images for PB/SBR blends are given in the SI (Figure S1). Note that small angle x-ray diffraction (SAXS) experiments do not provide further insights since scattering contributions due to nanoparticles, additives and rubber matrix do strongly superimpose for block copolymer based composites. Dynamic shear measurements for a PB-SBR block copolymer based composite containing up to 80 phr silica and the unfilled reference sample show as expected that there are two segmental (α) relaxation processes related to softening of the two polymeric phases (Figure 2). The PB domains do soften at about −83 ◦ C while the chosen SBR block softens at about −6 ◦ C if the peak maximum in the temperature-dependent loss modulus G00 (T ) measured at 10 rad/s is considered as criterion (SI, Figure S2). These values correspond relatively well to those of cross-linked PB and SBR rubbers with comparable microstructure. Note that there is also no significant shift in the α peak positions comparing the shear curves for the cross-linked PB-SBR block copolymer with those of the composites. This can be concluded

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G ' / P a

S ilic a -fille d P B

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Figure 2: Master curves for (a) storage modulus G’, (b) loss modulus G” and (c) loss tangent of vulcanized PB-SBR diblock copolymers containing different amounts of silica. The original isotherms are measured in the frequency range 0.1 - 100 rad/s. Reference temperature is -60 ◦ C and the individual isotherms are shifted horizontally in the entire temperature range.

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from the positions of the αP B and αSBR peaks in G00 for different composites in master curves (Figure 2b), isochrones (SI, Figure S2b) as well as a comparison of tanδ isotherms measured close to Tg,P B and Tg,SBR (SI, Figure S3). Shift factors aT taken from master curve constructions do also support that the temperature dependence of the α relaxation times τα = ωα−1 are basically independent of the filler content (SI, Figure S4). The trends in reinforcement and dissipation depending on filler content and temperature become clear if dynamic strain sweeps for block copolymer based composites are compared (Figure 3). As expected, a strong increase is observed in the storage modulus of filled rubbers at very small strain amplitudes (G00 ) with increasing filler content. This holds for all investigated temperatures (0, 25, 60 ◦ C) although the absolute G00 values are much smaller for higher temperatures. With increasing strain amplitude the G0 values decrease systematically for the investigated block copolymer based composites. This resembles features known as Payne effect 22 and can be understood as an indication of a filler network breakdown at large deformations. The dependence of the storage modulus values at very high amplitudes G0∞ on filler content and temperature is much less pronounced. Accordingly, the difference ∆G0 = G00 − G0∞ - representing the load carrying capacity of the filler network - decreases with decreasing filler content and increasing temperature like in other series of elastomer nanoparticle composites. 11,23 The temperature-dependence of ∆G0 has been interpreted as evidence for the visco-elastic nature of the filler network and the appearance of glassy rubber bridges connecting neighbored filler particles which soften gradually with increasing temperature (and decreasing measurement frequency). The solid lines in Figure 3a correspond to fits with the established Krauss equation 24

G0γ =

G00 − G0∞ + G0∞ 1 + ( γγc )2m

(1)

where G00 and G0∞ are the extrapolated moduli at 0 and ∞ strain amplitude, respectively. The parameter m is an exponent describing the decrease in G0 (γ) and γc is a characteristic strain amplitude where the filler network breaks. The m values at higher temperatures 8


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(25◦ C and 60◦ C) are about 0.6 like in many other rubber composites 11,14 while the γc values are commonly in the range 1% − 3% with a certain trend to decrease with increasing filler content. Note that strain sweeps performed on the unfilled block copolymer sample at 25◦ C and 60◦ C show expectedly no considerable change depending on strain amplitude (Figure 3). Peculiar behavior is only seen if an unfilled block sample is measured at 0◦ C. A Payne effect is in this case observed without filler since we are close to the glass temperature of the SBR phase (Tg,SBR ≈ −8◦ C). This clearly indicates that the SBR phase percolates in this self-assembled block copolymer rubber with lamellar morphology. This peculiarity is accompanied by slightly higher m values (about 0.75) from Kraus fits and an unusual decrease in the loss tangent at small deformations with increasing filler loading which is commonly not observed well above the glass temperature. Considering G00 (γ) data from strain sweeps, trends are seen which are also observed in other composites. Commonly the G000 values at small strain amplitudes are higher than those at extremely large strain amplitudes G00∞ while a peak is observed at intermediate strain amplitudes. The peak in G00 (γ) has been related already by Kraus to dissipation due to breaking of bridges in the filler network. The difference ∆G00 = G000 − G00∞ has been recently related to dissipation caused by oscillatory deformation of intact glassy rubber bridges in the filler network. 11 Since the number of intact glassy rubber bridges in the filler network decreases effectively with increasing strain amplitude, these contributions to dissipation show a behavior similar to that of G0 (γ). The step-height ∆G00 as well as the peak height decrease systematically with decreasing filler content and increasing temperature. The solid lines in Figure 3b are fits based on the equation 11,25

G00γ

2(G00m − G00∞ )( γγc )m G000 − G00∞ + + G00∞ = 1 + ( γγc )2m 1 + ( γγc )2m

(2)

describing G00 (γ) curves from strain sweeps usually quite well. G00∞ is the value of the loss modulus at very high amplitudes where the filler network is broken. A linear relation 9



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Figure 3: Strain sweeps for (a) storage modulus G’, (b) loss modulus G” and (c) loss tangent of vulcanized PB-SBR diblock copolymers containing different amounts of silica (cf. labels) measured at 0 ◦ (top), 25 ◦ C (middle), 60 ◦ (bottom) and an angular frequency of 10 rad/s.

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0,10 PB/SBR blend based composites 0,08 tan (60°C)

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PB-SBR block copolymer based composites

0,06

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Figure 4: Rolling resistance compound indicator (tan δ(60)◦ C) vs. reinforcement (G’(60◦ C)) for block copolymer based composites (full symbols) and blend based composites (open symbols).

between step height (G000 −G00∞ ) and peak height (G00m −G00∞ ) has been observed for composites with percolated filler network supporting the idea that both quantities are proportional to the number of glassy rubber bridges in the filler network in the undeformed state. 11 From the applied point of view the ratio of dissipation in the rolling resistance relevant range versus reinforcement is very important. This ratio should be small in order to get a good tire tread performance. In case of compound characterization, the storage modulus G00 (60◦ C) and the loss tangent tanδ(60◦ C) measured at 10 rad/s might be useful lab indicators allowing to estimate reinforcement and dissipation in the range which is relevant for the rolling resistance of tires. Both quantities are plotted against each other in Figure 4. It is clear that G00 (60◦ C) as well as tan δ(60◦ C) increase similarly with increasing filler content resulting in nearly linear dependence. Most interesting is, however, a comparison with corresponding results for a PB/SBR blend filled with identical amounts of silica particles. Main finding is that the compound indicators tanδ(60◦ C) for blend based composites (related to rolling resistance) are generally significantly higher than those of the block copolymer based composites for a given reinforcement (G00 (60◦ C)). The original strain sweep data for filled blends are given in Figure S5 of the SI. Qualitatively, all trends seen in the strain

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sweeps for the filled PB-SBR block copolymer are preserved and fits based on Eqs. (1) and (2) are still a good approximation. It is obvious that the storage modulus at low deformations is always significantly lower as compared to the corresponding values for block copolymer based composites containing the same amount of filler. Note that the elastomers used in these blend based composites are commercial grades having a similar microstructure like the blocks of the investigated PB-SBR copolymer but higher average molecular weights (for details see SI). Considering information about the filler dispersion in the investigated blend based composites as observed from AFM images (Figure S1) one can conclude that it is significantly more homogeneous than in block copolymer based composites for high filler contents (≥ 60phr) although the silica particles are also preferentially located in the SBR phase of the blend with low filler contents (20phr). We think that this difference might be a major reason for the improved mechanical performance of block copolymer based composites as shown in Figure 4. The combination of a high degree of localization of the silica particles in the SBR phase at application-relevant filler loadings (≥ 60 phr) with a nanoscale morphology of the block copolymer matrix may lead to a favorable filler network topology resulting in a high reinforcement combined with relatively low dissipation in the range relevant for the rolling resistance of tire treads. The key to understand these advantageous features might be the high concentration of silica in a continuous SBR phase probably connected with a lower percolation threshold in our block copolymer based composites as compared to classical counterparts. 26 Summarizing the results of this study one can conclude that PB-SBR block copolymer based composites with application relevant silica loadings (≥ 60 phr) show a very high degree of filler selectivity towards the SBR phase compared to blend based composites. The nearly exclusive location of silica particles in the SBR domains gives rise to a special filler network topology which is to a large extent influenced and controlled by the self-assembled morphology of the block copolymer. Most interesting is that this filler network topology can help to come to favorable mechanical properties in case of rubber composites for tire treads. Appli-

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cation relevant values for reinforcement (G00 (60◦ C)) and hardness can be achieved for lower filler loadings in block copolymer based composites as compared to other rubber matrices. With this reduction regarding the required filler contents dissipation values like tanδ(60◦ C) do also significantly drop down since such quantities are usually roughly proportional to the filler content. For the investigated block copolymer based composites a reduction in tanδ(60◦ C) of more than 25% is obtained at application relevant reinforcement values of about 10 MPa in comparison to blend PB-SBR based composites containing the same type of filler (Figure 4). This clearly indicates a significant potential of the presented approach for improving the performance of rubber composites for tire treads.

Acknowledgments. The authors thank Goodyear S.A. for allowing us to publish the data shown and the Fonds National de la Recherche in Luxembourg (FNR) for financial support (grant number 4797207).

ASSOCIATED CONTENT Supporting Information available: Polymer characteristics, additional shear data for PB-SBR block copolymer based composites as well as strain sweeps and AFM images for PB/SBR blends are provided.

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References (1) Bockstaller, M. R.; Mickiewicz, R. A.; Thomas, E. L. Block Copolymer Nanocomposites: Perspectives for Tailored Functional Materials. Advanced Materials 2005, 17, 1331–1349. (2) Lin, Y.; Boker, A.; He, J.; Sill, K.; Xiang, H.; Abetz, C.; Li, X.; Wang, J.; Emrick, T.; Long, S.; Wang, Q.; Balazs, A.; Russell, T. P. Self-Directed Self-Assembly of Nanoparticle/Copolymer Mixtures. Nature 2005, 55 – 59. (3) Haryono, A.; Binder, W. W. H. Controlled Arrangement of Nanoparticle Arrays in Block-Copolymer Domains. Small 2006, 2, 600–611. (4) Kim, B. J.; Bang, J.; Hawker, C. J.; Kramer, E. J. Effect of Areal Chain Density on the Location of Polymer-Modified Gold Nanoparticles in a Block Copolymer Template. Macromolecules 2006, 39, 4108–4114. (5) Vaia, R. A.; Maguire, J. F. Polymer Nanocomposites with Prescribed Morphology: Going beyond Nanoparticle-Filled Polymers. Chem. Mater. 2007, 19, 2736 – 2751. (6) da Silva, C. A.; Budde, H.; Menzel, M.; Wendler, U.; Bartke, M.; Weydert, M.; Beiner, M. Self-Assembled Structure and Relaxation Dynamics of Diblock Copolymers Made of Polybutadiene and Styrene/Butadiene Rubber. RSC Adv. 2016, 6, 50460– 50470. (7) Thielen, G. Filler Impact on Rubber Blend Miscibility. Kautsch. Gummi Kunstst. 2007, 60, 389–392. (8) da Silva, C. A.; Weydert, M.; Budde, H.; Wendler, U.; Menzel, M.; Bartke, M.; Beiner, M. Interrelation Between Morphology and Softening Behavior in Self-Assembled Poly(butadiene-block-(styrene-stat-butadiene) Copolymers. Rubber Chem. Technol. 2016, 89, 392–405. 14


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(9) Wang, M. J. Effect of Polymer-Filler and Filler-Filler Interactions on Dynamic Properties of Filled Vulcanizates. Rubber Chem. Technol. 1998, 71, 520–589. (10) Heinrich, G. The Dynamics of Tire Tread Compounds and Their Relationship to Wet Skid Behavior. Prog. Colloid Polym. Sci. 1992, 90, 16–26. (11) Nagaraja, S.; Mujtaba, A.; Beiner, M. Quantification of Different Contributions to Dissipation in Elastomer Nanoparticle Composites. Polymer 2017, 111, 48–52. (12) Heinrich, G.; Klüppel, M. Recent Advances in the Theory of Filler Networking in Elastomers. Adv. Polym. Sci. 2002, 160, 1–44. (13) Merabia, S.; Sotta, P.; Long, D. R. A Microscopic Model for the Reinforcement and the Nonlinear Behavior of Filled Elastomers and Thermoplastic Elastomers (Payne and Mullins Effects). Macromolecules 2008, 41, 8252–8266. (14) Mujtaba, A.; Keller, M.; Ilisch, S.; Radusch, H.-J.; Thurn-Albrecht, T.; Saalwächter, K.; Beiner, M. Mechanical Properties and Cross-Link Density of Styrene-Butadiene Model Composites Containing Fillers with Bimodal Particle Size Distribution. Macromolecules 2012, 45, 6504–6515. (15) Robertson, C. G.; Lin, C. J.; Bogoslovov, R. B.; Rackaitis, M.; Sashukhan, P.; Quinn, J. D.; Roland, C. M. Flocculation, Reinforcement, and Glass Transition Effects in Silica-Filled SBR. Rubber Chem. Technol. 2011, 84, 507–519. (16) Mujtaba, A.; Keller, M.; Ilisch, S.; Radusch, H.-J.; Thurn-Albrecht, T.; Saalwächter, K.; Beiner, M. Detection of Surface-Immobilized Components and Their Role in Viscoelastic Reinforcement of Rubber-Silica Nanocomposites. ACS Macro Lett 2014, 3, 481–485. (17) Heinrich, G.; Klüppel, M.; Vilgis, T. A. Reinforcement of Elastomers. Curr. Opin. Solid State Mater. Sci. 2002, 6, 195–203.

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(18) Otegui, J.; Miccio, L. A.; Arbe, A.; Schwartz, G. A.; Meyer, M.; Westermann, S. Determination of Filler Structure in Slilica-Filled SBR Compounds by Means of SAXS and AFM. Rubber Chem. Technol. 2015, 88, 690–710. (19) Ganter, M.; Brandsch, R.; Thomann, Y.; Malner, T.; Bar, G. Recent Progress in Atomic Force Microscopy of Elastomers, TPE Blockcopolymers and Blends. Kautsch. Gummi Kunstst. 52, 717–723. (20) Inai, M.; Aizawa, S.; Ito, M. Phase Control of BR/SBR Blends by Silica Particles. Soft Matter 2007, 3, 64–69. (21) Le, H.; Keller, M.; Hristov, M.; Ilisch, S.; Xuan, T.; Do, Q.; Pham, T.; Stockelhuber, K.; Heinrich, G.; Radusch, H. Selective Wetting and Localization of Silica in Binary and Ternary Blends Based on Styrene Butadiene Rubber, Butadiene Rubber, and Natural Rubber. Macromol. Mater. Eng. 2013, 298, 1085–1099. (22) Payne, A. R. The Dynamic Properties of Carbon Black-loaded Natural Rubber Vulcanizates. J. Appl. Polym. Sci. 1962, 6, 57–63. (23) Reinforcement factors G’(composite)/G’(rubber) as used in other cases would behave qualitatively similar. However, we do not favor this quantity for the description of filler network related contributions to reinforcement since there is no rubbery polymer incorporated directly in the filler network according to the physical picture used here. Further, we think that reinforcement factors do not reflect the additivity of reinforcement contributions caused by percolating filler network, occluded rubber and hydrodynamic reinforcement of the rubber matrix which is commonly assumed for rubber composites. (24) Kraus, G. Mechanical Losses Carbon-Black Filled Rubbers. Appl. Polym. Symp. 1984, 39, 75–92. (25) Ulmer, J. D. Strain Dependence of Dynamic Mechanical Properties of Carbon Black Filled Rubber Compounds. Rubber Chem. Technol. 1996, 69, 15–47. 16


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(26) Note that the morphology of the block copolymer matrix should have a strong influence on the storage modulus G’ as long as the filler is located exclusively in the SBR phase. Whether or not related effects in blend based composites are pronounced is much less clear and somehow more unlikely. In blends, the domains are usually much larger and significantly less defined. Long range order is missing in blends and the ’morphology’ of the rubber matrix is defined by the processing procedure. Moreover, the degree of filler selectivity is usually much lower. Without high selectivity of the filler particles for one rubber phase there is no reason for an increase in G’.

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Diblock Copolymer Based Composites for Tire Tread Applications with

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