Dynamic Moduli Mapping of Silica-Filled Styrene–Butadiene Rubber

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Dynamic Moduli Mapping of Silica-Filled Styrene−Butadiene Rubber Vulcanizate by Nanorheological Atomic Force Microscopy Eijun Ueda,†,‡ Xiaobin Liang,‡ Makiko Ito,‡ and Ken Nakajima*,‡ †

Zeon Corporation, Kanagawa 210-9507, Japan Department of Chemical Science and Engineering, Tokyo Institute of Technology, Tokyo 152-8552, Japan

‡

Macromolecules Downloaded from pubs.acs.org by UNIV OF EDINBURGH on 12/21/18. For personal use only.

S Supporting Information *

ABSTRACT: Atomic force microscopy (AFM) offers nanoscale mapping of materials’ properties. Especially, our modified AFM termed “nanorheological AFM” enables us to measure the accurate frequency-dependent storage and loss moduli and loss tangent over a sixth-order frequency range without any temperature control at nanoscale resolution. These dynamic properties obtained by nanorheological AFM can be compared with those using bulk dynamic mechanical analysis (DMA) measurements. In this paper, we applied this technique to silicafilled styrene−butadiene rubber (SBR) to investigate the nature of the interfacial rubber region existing between a rubber matrix and silica particles at different temperatures. The dynamics properties of the interfacial rubber region were different from those of matrix rubber regions. The master curve obtained by this technique perfectly coincided with that by bulk DMA. Furthermore, it was found that the behavior of bulk loss tangent could be predicted by the contributions from both matrix and interfacial rubber regions, which assures the importance of “interface control” often adopted in the tire industry to change tire performances. By examining the frequency-dependent change of loss modulus from the glass−rubber transition state to the glassy state, we also found for the first time that this transition does not occur uniformly but inhomogeneously depending on spatially heterogeneous polymer segmental dynamics.

1. INTRODUCTION

heterogeneous PNCs. Furthermore, there is no guarantee for TTS to be used in nanoscale measurements. To date, atomic force microscopy (AFM)7 has been widely used to investigate inhomogeneous structures such as rubbery materials, since AFM offers us the investigation of surfaces and interfaces of such materials. Especially, AFM enables us to visualize not only their structural information but also several mechanical property maps like Young’s modulus in nanoscale resolution. Successful measurements on some dynamic properties such as storage and loss moduli and tan δ was even reported up to now.8−15 For example, Ding et al. have reported how AFM techniques based on contact resonance (CR-AFM)8 can be used to measure the viscoelastic loss tangent of four homogeneous polymeric materials. The method would lead a meaningful consideration in obtaining tan δ data on homogeneous materials, while reliable data can be obtained only in narrow frequency range since the CR-AFM is operated at near-resonant frequency of the vibrating AFM cantilever in principle. Therefore, it was necessary to intentionally change the temperature for obtaining viscoelastic properties of specimens in their report. This method is not suitable for softer and adhesive materials such as rubber. As mentioned before, there is no guarantee for TTS to be applicable to

Recently, ecology tires have become increasingly popular to reduce greenhouse gas emissions. The reduction of their rolling resistance (RR) has a significant impact on fuel savings of vehicles. These tires are made from polymer nanocomposites (PNCs) which are often composed of styrene− butadiene rubber (SBR) and silica particles to improve tire performances such as RR efficiency and wet grip (WG) property. Among several dynamic properties of PNCs, loss tangent tan δ takes an important role as a good predictor of RR and WG for automobile tires.1−4 In PNCs, there exist highly heterogeneous and hierarchical structures with various length scales5 ranging from nanometers to micrometers and various frequency scales ranging from hertz to megahertz, affecting the demand characteristics for automobile tires. However, it is difficult to predict macroscopic dynamic properties corresponding to RR or WG only from the structural information such as nanoparticle distributions and some characteristic phase morphologies. For bulk specimens, as a general concept, frequency and temperature can be converted mutually based on the time−temperature superposition (TTS)6 principle so that one can have access to highfrequency responses from the specimens not by having such high-frequency measurements but just by cooling the specimens. However, TTS is an empirical rule, and more importantly it would not be so simple to apply TTS to © XXXX American Chemical Society

Received: October 22, 2018 Revised: December 14, 2018

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Macromolecules nanoscale mechanical measurements. One cannot expect that his experiments reproduce the situation of real contact between roads and a tire, even by adjusting temperature and frequency. We need a tool to obtain viscoelastic behavior of polymeric materials only by adjusting frequency in a wide range. Nanofillers such as carbon black (CB) or silica particles with different diameters are used to improve the mechanical properties in the automobile industry. There are many controversies about the reinforced mechanism using nanofillers.16,17 Many rubber scientists to date concluded that there may be a glassy rubber phase called “bound rubber” around nanoparticles such as silica particles or CB aggregates in filledrubber compounds.18−21 As previously studied by pulsed NMR, the relaxation time of the bound rubber region was shorter than that of matrix rubber region, indicating the restricted segmental motion.22−24 However, pulsed NMR only gives information about relaxation time averaged over bulk specimens and therefore lacks spatial information and cannot offer the direct evidence of bound rubber. The method cannot tell us where the bound rubber exists. These restricted rubber regions may show a harder response than the other matrix rubber; in other words, the interfacial rubber regions around silica particles could indicate a higher glass−rubber transition temperature (Tg) than the matrix. In our previous studies, the static, yet quantitative mechanical property mapping by nanopalpation AFM was employed on rubbery materials such as a CB-filled isoprene rubber (IR) and could visualize harder interfacial rubber regions surrounding CB aggregates,25 where the analysis was based on the Johnson−Kendall−Roberts (JKR) contact mechanics model. Furthermore, aiming at dynamic mechanical property mapping, we have developed the nanorheological AFM26 with a wide frequency range from 1 to 104 Hz in SBR/ IR blends at room temperature. The nanorheological AFM combines static Young’s modulus determination, which calculates the contact area between a specimen surface and an indenting AFM probe, and dynamically oscillating excitation to perform storage and loss moduli measurement very quantitatively. Because nanorheological AFM does not rely on the resonance of cantilever, which is a main difference with CR-AFM, a frequency sweep in wide range is possible. In the case of the SBR/IR blend specimen, we could conclude that the SBR phase reached Tg and behaved as glassy phase at higher frequencies. Therefore, it becomes easier to predict the segmental dynamics of polymer chains and thus mechanical properties at nanoscale by this method. In this article, we will introduce a modified nanorhelogical AFM method to measure the viscoelastic properties of polymeric materials, especially rubbers, over a wide frequency range and a wide temperature range. In particular, we investigated frequency-dependent dynamic moduli and loss tangent quantitatively at a matrix rubber region and an interfacial rubber region of silica-filled SBR in detail with nanoscale resolution.

Table 1. Compound Recipe of SBR Vulcanizate

a

constituent

phra

styrene−butatiene rubber silica (Zeosil 1165MP) bis[3-(triethoxysilyl)propyl]tetrasulfide N-(1,3-Dimethylbutyl)-N′-phenylphenylenediamine stearic acid N-cyclohexyl-2-benzothiazolesulfenamide 1,3-diphenylguanidine sulfur

100 15 1.2 2 2 2.5 2 2

phr is parts by weight per 100 parts by weight of rubber.

with a glass and a diamond knife prior to AFM measurements and placed on the flat Si wafer substrate attached on piezoelectric actuator. Nanorheological AFM Measurement. A hard diamond-like carbon-coated AFM cantilever (tap300DLC, BudgetSensors, Bulgaria) with spring constant of 40.2 N/m (calibrated) and resonant frequency of 264 kHz (measured) was used for this measurement. The radius of curvature of the AFM probe was also measured to be ∼100 nm, a little bit larger aiming at a stable measurement. A modified commercial AFM (NanoScope V with FastScan Icon, Bruker AXS, USA) was used. In this study, to measure the viscoelastic properties of polymeric materials over a wide frequency range, our method includes a modified AFM instrument, hereafter termed nanorheological AFM, which has a tiny piezoelectric actuator placed between the specimen and AFM sample stage. The specimen is oscillated by this actuator with a frequency between 1 Hz and 20 kHz. The measurement temperature was controlled from −10 to 60 °C with thermal applications controller (TAC, Bruker AXS, USA) under a nitrogen atmosphere. When the cantilever deflection value reaches the preset force of 50 nN during each force−distance curve measurement (a typical curve is shown in Figure S1 of the Supporting Information), the tiny actuator starts to oscillate, and the AFM probe is maintained in contact with the sample surface until the preset frequency sweep ranging from ∼1 Hz to 20 kHz is completed. The oscillation amplitude of the actuator was 5.0 nm in this study. Because our nanorheological AFM calculates static modulus at each point based on JKR analysis,25−27 the actual contact area is obtained. Eventually one can convert dynamic stiffness (S′: storage stiffness; S″: loss stiffness), which is the direct measurand of nanorheological AFM, into dynamic moduli (E′: storage stiffness; E″: loss stiffness) as follows: E′ =

(1 − v 2)S′ , 2at C(at )

E″ =

(1 − v 2)S″ 2at C(at )

(1)

where ν is Poisson’s ratio and C(at) is C(at ) =

1 − (a0 /at )3/2 1 − (1/6)(a0 /at )3/2

(2)

a0 in this equation is the contact radius at the balance point, where adhesive force and elastic restoring force balance each other. The oscillation with the amplitude of a few nanometers (5.0 nm in this experiment) is executed at a = at (usually the repulsive-force region). The loss tangent is determined to be tan δ ≡

E″ S″ = E′ S′

(3) 26,27

For more details, refer to our previous works. Bulk DMA Measurement. The viscoelastic properties for a bulk specimen were obtained by a temperature (T)−frequency ( f) dispersion measurement using dynamic mechanical analysis (DMA, ARES-G2, TA Instruments, USA). The raw data were obtained by the shear mode under a constant strain amplitude in a frequency range of 0.05 to 50 Hz over a temperature range of −50 to +45 °C. The shear modulus (G′, G″) can be converted into the tensile modulus (E′, E″) using the following equations: G′ = E′/2(1 + ν) and G″ = E″/2(1 +

2. EXPERIMENTAL SECTION Materials. SBR was prepared with living anion polymerization and has a molecular weight of almost 370000. Table 1 shows the recipe of the SBR vulcanizate. The Tg of the filled vulcanizate measured by differential scanning calorimetry (DSC7000X, Hitachi, Japan) was −20.3 °C. The specimens were cut into a 1.5 μm thick thin film using an ultramicrotome (UC7 Leica Microsystems, Germany) at −100 °C B

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Figure 1. Static elastic modulus maps obtained by AFM controlled at (a) 26 °C and (b) 11 °C. A scale bar indicates 200 nm length. Both maps were captured at 1 Hz scan rate. The modulus histograms (c, d) were obtained from the original maps (a, b). ν). Here, Poisson’s ratio, ν, is assumed to be 0.5 for rubber materials. The raw data were shifted to obtain master curve, i.e., the relationship between reduced frequency, aT f, E′, E″, and tan δ using the Williams− Landel−Ferry (WLF) equation.6 The shift factor (aT) is calculated as

log aT = −

C1(T − Tref ) C2 + T − Tref

modulus values are ranging from 3σ to 7σ. The region larger than Epeak + 7σ corresponds to the “filler region”, which was also confirmed by a maximum adhesion force map obtained simultaneously with the elastic modulus map. By this procedure, the region-of-interest (ROI) image of each region was made, which is shown in Figure S2. Figure 1b indicates the static modulus map at 11 °C, with EJKR_matrix of 6.60 ± 0.15 MPa and EJKR_interface of 13.8 ± 1.66 MPa. Its elastic modulus histogram is shown in Figure 1d. It is clear that the elastic modulus obtained at 11 °C has a higher and wider distribution than that of 26 °C. This result means that both matrix and interfacial rubber regions become harder as the measured temperature approaches the Tg of SBR. The contrast between the vulcanized SBR matrix region and the interfacial rubber region near silica particles indicates that the mechanical properties of these two regions differ significantly to each other, with showing inhomogeneous dispersion of silica particles. In addition to the “bound rubber”, which are also referred to glassy hard (GH) layer, a transition layer between the tightly bound layer and the bulk matrix, known as a loosely bound layer, often called a sticky hard (SH) layer, has been also suggested to exist by Fukahori.28 He reported the thickness of the SH layer is in a range of a few tens of nanometers. The thickness of the interfacial rubber region in Figure 1 was about 10−50 nm and much larger than that of reported bound rubber layer (∼2 nm) and rather similar to the SH layer. A similar result was also reported for the carbon black-filled isoprene rubber, which discussed the mechanical interface measured by AFM.25 Next, the novel findings obtained by nanorheological AFM are shown in Figure 2 corresponding to E′, E″ and tan δ maps at 26 °C. The data at 11 °C are given in Figure S3. A series of dynamic mechanical maps at each frequency were captured simultaneously with the static measurement shown in Figure

(4)

where we adopted the universal constant (C1 = 8.86, C2 = 101.6). The reference temperature Tref was set to be 29.2 °C to obtain the best master curve.

3. RESULTS AND DISCUSSION Nanorheological AFM was applied to silica-filled SBR vulcanizate over wide frequency and temperature. In our method, nanorheological mapping data were taken as 64 pixels × 64 pixels for an imaging area of 1.0 μm × 1.0 μm square region, where each pixel contained the data measured at 11 and 13 different oscillation frequencies at 26 and 11 °C, respectively. The static scan rate was 1.0 Hz. First of all, the static elastic modulus maps measured by nanopalpation AFM27 controlled at 26 and 11 °C are shown in Figures 1a and 1b, respectively. Figure 1a is the static modulus map based on JKR contact mechanics analysis with static elastic modulus in matrix rubber region, EJKR_matrix of 3.56 ± 0.04 MPa, and that in interfacial rubber region, EJKR_interface of 5.03 ± 0.35 MPa at 26 °C. The definitions of the matrix rubber region and the interfacial rubber region are described as follows: the peak modulus value, Epeak, and its standard deviation, σ, of static elastic modulus (EJKR) distribution are calculated by fitting the modulus histogram in Figure 1c, obtained from the original map in Figure 1a, against a single Gaussian function. Then, we define the region corresponding to Epeak ± 1.5σ as the “matrix rubber region”. Next, the interfacial rubber region is defined as a region of which C

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Figure 2. Some nanorheological maps that have been extracted from original 33 (11 × 3) image sets. The scan size is 1.0 μm. The measured temperature is 26 °C. a, b, and c correspond to E′, E″, and tan δ, respectively. The measurement frequencies were 1 = 30 Hz, 2 = 100 Hz, 3 = 300 Hz, 4 = 2 kHz, and 5 = 5 kHz.

1a. Although over ∼10 h was required to capture them, the images in Figure 2 were obtained simultaneously and therefore from the same location, which is the advantage of this method. By comparing the matrix rubber regions of E′ and E″ maps at frequency of 30 Hz, E′ is larger than E″, typical rubbery behavior. As frequency increased, however, E″ of matrix rubber region became first comparable to E′ at 300 Hz and 2 kHz and

became larger than E′ at 5 kHz. In tan δ maps, the contrast gradually became clearer at higher frequencies. The frequency dependence of dynamic behavior such as E′ and tan δ is shown in Figure 3. These E′ and tan δ were calculated from nanorheological maps of all measured frequencies at 26 and 11 °C. As the frequency increased from 10 Hz to 20 kHz at 26 °C, tan δ statistically averaged D

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To fully elucidate the difference in segmental dynamics of polymer chains in SBR matrix and the interfacial rubber regions near silica particles in detail, we applied the abovementioned ROI images (see Figure S2) to divide tan δ maps into interfacial and matrix rubber regions. We do not think tan δ value on silica particles is correct because the silica particles are floating in rubber matrix, and the silica itself is never deformed. Thus, we exclude any effect of the filler regions in the following discussions. Figure 5 shows the frequency dependence of tan δ at each region obtained by nanorheological AFM controlled at 26 and 11 °C. Three points (a), (b), and (c) inserted in the graph correspond to the data with oscillation frequency of 10 Hz, 10 kHz at 26 °C, and 5 kHz at 11 °C, respectively. The corresponding tan δ histograms are also shown in Figure 5. First of all, focusing on the shapes of tan δ histogram, it is evident that the segmental dynamics of polymer chains had a wide distribution. In general, only one “averaged” dynamic behavior is obtained from one frequency by bulk DMA; however, our nanorheological AFM enable us to visualize tan δ dispersion maps on the surface of polymeric materials. The dynamic information from the matrix rubber regions and the interfacial rubber regions can be grasped separately and quantitatively. The histogram (a) for 10 Hz (aT f = 22.4 Hz) in Figure 5 indicates tan δ of the interface region is almost similar to that of the matrix, and their distributions are narrower. Though at elevated frequencies, the tan δ distribution of both regions became wider as shown in histogram (b). The degree of the distribution is also visible as the error bar in the graph. Furthermore, the difference of the tan δ peak value became evident. At all middle-range frequencies, this specimen showed the lower tan δ for the interfacial rubber region, possibly due to the restricted dynamics for the bound polymers at the interface. This tendency continued until about aT f = 18 kHz, followed by the decrease of tan δ in both the matrix and interfacial rubber regions. The rate of the decrease at the matrix was larger, and once again tan δ became similar to that of the interface at around aT f = 200 kHz. At such higher frequencies, this vulcanized sample had undergone the glass transition, and thus the polymer chains were in the glassy state, assuring no noticeable difference in each region. The peak frequency of tan δ for the interfacial rubber region slightly shifted toward the lower frequency, which might mean the Tg of the interface shifted to higher temperature; in other words, the segmental dynamics at the interface were restricted. This result may imply that the chain mobility in interfacial rubber regions was restricted by the chemical bonds formed between silane coupling agent and polymer chain.29 We will discuss this peak shift issue later again since the use of tan δ is not adequate in discussing the relaxation process of polymeric materials in general. As discussed above in Figure 4, tan δ averaged over the whole image coincided with that of the bulk specimen. However, as can be seen in Figure 5, tan δ of matrix rubber region shifted to higher values and that of interfacial rubber region to lower values. This observation may infer that the “additive law” holds in discussing the overall tan δ value. However, the “additive law” should not be discussed for tan δ, but for E′ and E″. The additivity never holds for the tan δ value, by definition. The detailed discussion is made in Figure S4 and its figure caption in the Supporting Information where we could not perfectly confirm the “additive law” in E′ and E″. Nevertheless, as an experimental fact, the tan δ of bulk DMA was seem to be sandwiched between that of the matrix and

Figure 3. Frequency dependence of the dynamic behavior, E′ and tan δ measured by nanorheological AFM measured at 26 and 11 °C.

over the all areas changed from 0.39 to 1.7 against the change of the frequency from 10 Hz to 7 kHz, followed by a decrease from 1.7 to 1.5 between 7 and 20 kHz. On the other hand, in the case of 11 °C, as the frequency increased from 30 Hz to 10 kHz, the value of tan δ increased from 1.3 to 1.6 between 30 and 200 Hz and then decreased from 1.6 to 0.44 between 200 Hz and 10 kHz. An interesting observation is that we can see the peaks of tan δ appear at 200 Hz and 7 kHz for different temperatures of 11 and 26 °C, respectively. Next, the master curve was obtained from the original data in Figure 3 and is shown in Figure 4. To obtain this curve, we

Figure 4. Master curve obtained by DMA and nanorheological AFM. The reference temperature, Tref, was identical for both measurement and was set to 29.2 °C.

used Tref = 29.2 °C to calculate the shift factor, aT, at 26 and 11 °C based on eq 4. The bulk DMA master curve is also superimposed in Figure 4. Interestingly, E′ and tan δ curves at 11 and 26 °C were connected to each other smoothly, and master curves for bulk and nanoscale perfectly coincided with each other. Consequently, nanorheological AFM has capability of predicting the macroscopic mechanical properties quantitatively. E

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Figure 5. Frequency dependence of tan δ separated into matrix and interfacial rubber regions. Open symbols indicate matrix rubber region and closed symbols interfacial rubber region. The bulk tan δ is also superimposed. Loss tangent histograms (a), (b), and (c) are corresponding to the data with oscillation frequency of 10 Hz, 10 kHz at 26 °C, and that of 5 kHz at 11 °C, respectively.

interfacial rubber regions. This would be a good message to rubber industry since R&D people in this field empirically know the importance of “interfacial control” to change bulk properties, especially loss tangent of tire materials which influence RR or WG performance. However, we would like to state that the industrial people should use loss modulus in place of tan δ. Figure 6 shows the frequency dependence of the standard deviation of tan δ, σtanδ, which was calculated by single Gaussian fitting for tan δ histogram obtained at 11 °C. As frequency increases, σtanδ, that is, the spatial distribution of each region, gradually increases at lower frequencies. There is a peak at about aT f = 18 kHz which corresponds to the tan δ peak frequency. This observation indicates that spatial heterogeneity is enhanced near glass-transition region. Moreover, the σtanδ of interface is larger than σtanδ of the matrix. We will discuss the issue of spatial heterogeneity later in more detail. It was mentioned in the previous paragraphs that Tg at the interface might shift to higher temperature by looking at the shift of the peak frequency of tan δ. However, this is not a correct consideration. We need to check the change in E″. Figure 7 shows the frequency dependence of E″ at 11 °C. The closed symbols are the data averaged over ROI areas (circle for matrix and square for interface). The interfacial E″ had higher values than that of the matrix at all frequencies, though the peak frequency, which is sought to be corresponding to Tg, almost coincided with each other at about aT f = 350 kHz. This

Figure 6. Relationship between the standard deviation of tan δ and the reduced frequency at 11 °C. Open and closed symbols correspond to matrix rubber region and interfacial rubber region, respectively.

means Tg of the interface and matrix rubber regions are identical. Furthermore, at frequencies higher than 350 kHz, σtanδ in both interface and matrix is almost the same and small as shown in Figure 6, indicating spatial heterogeneity was minimized. From these observations we can infer that the state of the specimen turns from glass-transition state into glassy F

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loss tangent histograms. However, the powerfulness of nanorheological AFM is found in its spatially resolved images. Undoubtedly, spatial heterogeneity should account for the wider distribution. The frequency-dependent changes of E″ at two selected points are superimposed in Figure 7. To indicate the location of these points, the loss modulus maps, masked by the corresponding ROI image, for both interfacial rubber regions and matrix rubber regions are shown in Figure 8. At the point indicated by an arrow in Figure 8a where ROI was interfacial rubber regions, the peak frequency of the particular point (open square symbol in Figure 7), which is located very near a silica particle among the interfacial rubber region, was shifted to lower frequency than that of the “average” interfacial rubber region or that of the average matrix region, indicating the segmental mobility there was the lowest possibly due to the chemical bonds formed between silane coupling agent and polymer chain.29 Conversely, the peak frequency of a certain point among the matrix rubber region (indicated by an arrow in Figure 8b and open square in Figure 7) was shifted to higher than that of the “average” matrix region. The point would have less interaction with silica particles. In summary, in silica-filled SBR vulcanizate, we found for the first time that the glass− rubber transition does not occur uniformly but inhomogeneously depending on spatially heterogeneous polymer segmental dynamics.

Figure 7. Frequency dependence of loss moduli at 11 °C. The modulus peak of both matrix (circle sympols) and interfacial (square symbols) rubber regions appeared at about aT f = 350 kHz. Closed and open symbols correspond to the value averaged over ROI area and that of a single point, respectively.

state and the segmental dynamics of each region start to freeze at around this frequency. The above discussions were basically based on the average or the standard deviation obtained from dynamic modulus or

4. CONCLUSIONS A unique modified AFM-based nanorheological analysis was conducted to investigate the nature of the interfacial rubber

Figure 8. Loss modulus maps, masked by the corresponding ROI image, for both (a) interfacial rubber regions and (b) matrix rubber regions at aT f = 18 and 350 kHz. The scale bar is 200 nm. The arrows indicate the locations of the points inside the interfacial rubber region and the matrix region used in Figure 7, respectively. G

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(6) Williams, M. L.; Landel, R. F.; Ferry, J. D. The Temperature Dependence of Relaxation Mechanisms in Amorphous Polymers and Other Glass-forming Liquids. J. Am. Chem. Soc. 1955, 77, 3701−3707. (7) Binnig, G.; Quate, C. F.; Gerber, Ch. Atomic Force Microscope. Phys. Rev. Lett. 1986, 56, 930. (8) Hurley, D. C.; Campbell, S. E.; Killgore, J. P.; Cox, L. M.; Ding, Y. Measurement of Viscoelastic Loss Tangent with Contact Resonance Modes of Atomic Force Microscopy. Macromolecules 2013, 46, 9396−9402. (9) Yuya, P. A.; Hurley, D. C.; Turner, J. A. Contact-resonance atomic force microscopy for viscoelasticity. J. Appl. Phys. 2008, 104, No. 074916. (10) Yuya, P. A.; Hurley, D. C.; Turner, J. A. Relationship between Q-factor and sample damping for contact resonance atomic force microscope measurement of viscoelastic properties. J. Appl. Phys. 2011, 109, 113528. (11) Killgore, J. P.; Yablon, D. G.; Tsou, A. H.; Gannepalli, A.; Yuya, P. A.; Turner, J. A.; Proksch, R.; Hurley, D. C. Viscoelastic Property Mapping with Contact Resonance Force Microscopy. Langmuir 2011, 27, 13983−13987. (12) Yablon, Da. G.; Gannepalli, A.; Proksch, R.; Killgore, J.; Hurley, D. C.; Grabowski, J.; Tsou, A. H. Quantitative Viscoelastic Mapping of Polyolefin Blends with Contact Resonance Atomic Force Microscopy. Macromolecules 2012, 45, 4363−4370. (13) Noda, K.; Tsuji, M.; Takahara, A.; Kajiyama, T. Aggregation Structure and molecular motion of (glass-fiber/matrix nylon66) interface in short glass-fiber reinforced nylon 66 composites. Polymer 2002, 43, 4055−4062. (14) Satomi, N.; Tanaka, K.; Takahara, A.; Kajiyama, T. Effect of Internal Bulk Phase on Surface Viscoelastic Properties by Scanning Probe Microscopy. Macromolecules 2001, 34, 6420−6423. (15) Radmacher, M.; Tillmann, R. W.; Gaub, H. E. Imaging viscoelasticity by force modulation with the atomic force microscope. Biophys. J. 1993, 64, 735−742. (16) Robertson, C. G.; Lin, C. J.; Rackaitis, M.; Roland, C. M. Influence of Particle Size and Polymer-Filler Coupling on Viscoelastic Glass Transition of Particle-Reinforced Polymers. Macromolecules 2008, 41, 2727−2731. (17) Huang, M.; Tunnicliffe, L. B.; Thomas, A. G.; Busfield, J. J. C. The glass transition, segmental relaxations and viscoelastic behaviour of particulate-reinforced natural rubber. Eur. Polym. J. 2015, 67, 232− 241. (18) Smit, P. P. A. The glass transition in carbon black reinforced rubber. Rheol. Acta 1966, 5, 277−283. (19) Kraus, G.; Gruver, J. T. Thermal expansion, free volume, and molecular mobility in a carbon black-filled elastomer. J. Polym. Sci. Polym. Phys. 1970, 8, 571−581. (20) Wang, M.-J. Effect of Polymer-Filler and Filler-Filler Interactions on Dynamic Properties of Filled Vulcanizates. Rubber Chem. Technol. 1998, 71, 520−589. (21) Mujtaba, A.; Keller, M.; Ilisch, S.; Radusch, H.-J.; Beiner, M.; Thurn-Albrecht, T.; Saalwächter, K. Detection of Surface-Immobilized Components and Their Role in Viscoelastic Reinforcement of Rubber−Silica Nanocomposites. ACS Macro Lett. 2014, 3, 481−485. (22) Nishi, T. Effect of solvent and carbon black species on the rubber−carbon black interactions studied by pulsed NMR. J. Polym. Sci., Polym. Phys. Ed. 1974, 12, 685. (23) Lüchow, H.; Breier, E.; Gronski, W. Characterization of Polymer Adsorption on Disordered Filler Surfaces by Transversal 1H NMR Relaxation. Rubber Chem. Technol. 1997, 70, 747−758. (24) Litvinov, V. M.; Steeman, P. A. M. EPDM−Carbon Black Interactions and the Reinforcement Mechanisms, As Studied by LowResolution 1H NMR. Macromolecules 1999, 32, 8476−8490. (25) Nakajima, K.; Ito, M.; Wang, D.; Liu, H.; Nguyen, H. K.; Liang, X.; Kumagai, A.; Fujinami, S. Nano-palpation AFM and its quantitative mechanical property mapping. Microscopy 2014, 63, 193−208.

region existing between a rubber matrix and silica particles at different temperatures. E′, E″, and tan δ obtained by nanorheological AFM nicely corresponded to the master curves obtained by bulk DMA. Furthermore, it was found that the bulk behavior could be explained to some extent by the contributions from both matrix and interfacial rubber regions, which assure the importance of “interface control” empirically performed in rubber industry, especially the tire industry. By examining the frequency-dependent change of E″ from glass− rubber transition state to glassy state, we also found for the first time that this transition does not occur uniformly but inhomogeneously depending on spatially heterogeneous polymer segmental dynamics. In summary, nanorheological AFM enables us to discuss bulk viscoelastic properties of rubber composites by analyzing nanoscale viscoelastic properties, taking into account the contributions from both matrix and interfacial rubber regions.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b02258. Figure S1: a typical force (f)−deformation (d) curve on a rubber matrix region; Figure S2: region-of-interest (ROI) images obtained from Figure 1 in the main text; Figure S3: some nanorheological maps that have been extracted from original 39 (13 × 3) image sets; Figure S4: frequency dependence of (a) storage and (b) loss moduli to check the additive law in these quantities; Figure S5: a typical Lissajous figure obtained at a rubbery point with 10 kHz frequency and 2 nm sample deformation (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Xiaobin Liang: 0000-0003-2497-2085 Ken Nakajima: 0000-0001-7495-0445 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The work is partly supported by Impulsing Paradigm Change through Disruptive Technologies (ImPACT) Program, JST, Cabinet Office, Government of Japan.

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REFERENCES

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DOI: 10.1021/acs.macromol.8b02258 Macromolecules XXXX, XXX, XXX−XXX