Direct Mapping of Nanoscale Viscoelastic Dynamics at Nanofiller

Aug 2, 2018 - By means of loss tangent imaging in amplitude-modulated atomic force microscopy, we report a direct mapping of the nanoscale viscoelasti...
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Direct Mapping of Nanoscale Viscoelastic Dynamics at Nanofiller/ Polymer Interfaces Hung K. Nguyen,† Xiaobin Liang,‡ Makiko Ito,‡ and Ken Nakajima*,‡ †

Department of Applied Chemistry, Kyushu University, Fukuoka 819-0395, Japan Department of Chemical Science and Engineering, Tokyo Institute of Technology, Tokyo 152-8552, Japan

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‡

ABSTRACT: The mobility gradient of polymer chains near a solid interface plays an important role in explaining a complex behavior of relaxation dynamics and thermodynamic properties observed for various polymer nanocomposites. However, it still remains a major challenge to directly visualize this gradient at nanoscale. By means of loss tangent imaging in amplitude-modulated atomic force microscopy, we report a direct mapping of the nanoscale viscoelastic dynamics of interfacial polymers in two model nanocomposites including isoprene and styrene−butadiene rubbers filled with carbon black nanofillers. The loss tangent images clearly show a gradual slowdown of the viscoelastic dynamics of interfacial polymers from the bulk to the glassy behavior of a tightly bound layer over a length scale of ∼30 nm. These results can therefore provide important insights into understanding the longexisting question about the dynamic behavior at the nanofiller/polymer interfaces.

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INTRODUCTION The addition of nanofillers to a polymer matrix can substantially improve multiple macroscopic properties of the composite material. This improvement is largely attributed to the formation of an interfacial polymer region (IPR) in the proximity of the nanofiller surface with structural properties different from the bulk behavior.1−4 This IPR is also known as the dynamic interfacial region,5−7 implying that the segmental mobility and many associated quantities, such as the glass transition temperature (Tg), elastic modulus, and viscoelastic and dielectric responses, of interfacial polymers can be significantly altered from the bulk properties. A comprehensive understanding of the structural and dynamic changes within an IPR is therefore critical in tailoring multiple properties of soft materials filled with nanofillers.2−4,8 Consequently, a large number of experimental methods, theoretical models, and simulations have been devoted to describing the structural and dynamic behavior of interfacial polymers under the interaction with a solid surface over the past few decades.1,9−22 Theoretical models and simulations of interfacial polymers have generally suggested that the polymer chains in direct contact with a solid surface can form a tightly glassy bound layer of a few nanometers, which is connected with the bulk region through a transition interfacial layer, also known as loosely bound layer.5,21−25 The thickness of the loosely bound layer can vary from a few to tens of nanometers, depending on many factors, such as polymer molecular weight, geometry and size of nanofillers, and surface nature of the nanofiller/ substrate. The segmental dynamics of the tightly bound layer can be much slower than the bulk dynamics, whereas the mobility of polymer chains in the loosely bound layer should change from the bulk behavior to that of the tightly bound © XXXX American Chemical Society

layer. The idea about this mobility gradient of the IPR is critical to explain the complex behavior in the temperature, frequency, and strain dependence of mechanical, dielectric, and thermodynamic properties observed for various polymer nanocomposites.5,6,21,22,26,27 However, it still remains experimentally challenging to directly observe a dynamic IPR in real space. Many experimental efforts have been made based on an atomic force microscopy (AFM) to directly visualize the presence of an IPR surrounding nanofillers in various polymer nanocomposites.10−14,17,19 By using probe indentation AFM, modulus maps of various nanofilled polymers have been reported, which showed the existence of an IPR between the nanofiller and neat polymer matrix.12,13,19 The average modulus of the IPR was found to increase by a factor of ∼2 for glassy and ∼10 for rubbery polymers in comparison with that of the neat polymers. A remarkably clearer visualization of the IPR has been generally obtained using a phase imaging in amplitude-modulated (AM) AFM,10,13,17 a well-known method for determining the nanoscale variations in morphology, adhesion, dissipation, and composition of various complex interfaces.28−33 Relying on the phase imaging, a length scale of ∼20 nm for IPRs in several nanofilled polymers has been reported.13,17 Unfortunately, there still remain some inherent limitations in these AFM-based methods to clearly visualize the dynamic behavior of an IPR. The probe indentation AFM can only provide the information about the average static modulus of an IPR. Moreover, because of the requirement for a large Received: June 4, 2018 Revised: July 18, 2018

A

DOI: 10.1021/acs.macromol.8b01185 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Table 1. Sample Characteristics sample/phr

SBR

CB/SBR CB/IR

100

IR

CB

steric acid

ZnO

accelerator

sulfur

Tg/°C

100

5 10

1 1

2.5 5.0

1 1

1.5 2.0

−5 −65

where A and A0 are the amplitudes of the cantilever oscillation with and without the interaction with the sample, respectively, and φ is the phase shift of the cantilever oscillation from the external excitation signal. It is important to mention that unlike the dissipated process in the DMA method that is solely attributed to the viscoelastic nature of the sample, the Ploss in the AM-AFM might exhibit other contributions from the adhesive and capillary interactions between the tip and the sample surface.37,40 These contributions generally give rise to an overestimation of the AM-AFM tan δ in comparison with the DMA tan δ. These effects could, however, be significantly reduced by performing the experiment in the net repulsive regime and surrounded with a liquid/inert gas environment.37,40−42 AFM Measurements. AFM measurements were performed using a NanoScope V with the MultiMode 8 system (Bruker, USA) at room temperature (Troom ∼ 25 °C). Silicon cantilevers (Tap300Al-G) with a nominal tip radius of 5−10 nm were purchased from BudgetSensors. The resonant frequency f 0, spring constant k (calibrated), and quality factor Q of these cantilevers were ∼300 kHz, 40 N/m, and 500, respectively. A high value of Q (>100) was necessary for loss tangent measurements to eliminate the effect of higher harmonics of the cantilever.36 The actual tip radius was determined to be ∼7 nm with a so-called blind reconstruction method using a probe characterizer (TipCheck, Aurora NanoDevices, Canada).43 During each AM-AFM scanning, the topography, A, and φ images were obtained simultaneously, which represent the morphology of the sample surface, the change in the instantaneous cantilever amplitude at each measuring point, and the phase lag of the cantilever oscillation from the driving force. The latter two images were used to obtain tan δ image using eq 2. To reduce the effects of the adhesive and capillary forces on the tan δ imaging, which were mainly caused by an ultrathin water layer on the polymer surface,40 all measurements were performed in a N2 atmosphere using a gas flow rate of ∼5 L/min. By this way, the overestimation of the AM-AFM tan δ on neat SBR and IR regions with respect to those measured by a standard DMA was significantly reduced as shown in Figure 1.37

number of data per each force curve, this method could not provide a high-resolution mapping of an IPR.34,35 On the other hand, although the phase measurement has been proved capable of providing high-resolution mappings of an IPR, it is still difficult to correlate the phase shift to a specific property of the material.29,36 In this study, we employ a recently developed approach, namely loss tangent (tan δ) imaging, based on the AM-AFM method to directly visualize the viscoelastic dynamics of nanoscale IPRs at high spatial resolution.36,37 AM-AFM tan δ imaging results obtained here on two model nanocomposites including styrene−butadiene rubber (SBR) and isoprene rubber (IR) filled with carbon black (CB) nanofillers clearly show the existence of an IPR near CB interface with viscoelastic dynamics distinct from the bulk behavior. More importantly, the viscoelastic mobility along the interfacial region is directly evidenced to gradually slow down from the bulk behavior to that of a glassy layer bound to CB nanofillers.

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EXPERIMENTAL SECTION

Materials. Two types of nanofilled rubbers including HAF-grade CB-filled SBR vulcanizates and IR vulcanizates, which were kindly provided by Bridgestone Corp., were used as model specimens.38 These rubbers were thoroughly washed in different solvents to remove unreacted components before experimental characterization. The fraction of CB loading, sulfur content, vulcanization conditions, and Tg of these specimens are listed in Table 1. Tg values were measured by differential scanning calorimetry (Q200, TA Instruments). The specimens were cryo-faced using an ultramicrotome (UC6, Leica Microsystems, Germany) at −100 °C with a glass and a diamond knife prior to the AFM measurements. AM-AFM tan δ Imaging. In a standard technique, such as dynamic mechanical analysis (DMA), the viscoelastic dynamics of a polymeric material can be effectively described using tan δ measurements over a broad frequency/temperature range.39 This tan δ quantity is defined as the ratio of the loss energy to the stored energy at each cycle of perturbing the sample by a mechanically oscillating force. Taking into account this idea from the macroscopic measurements, Proksch and Yablon have recently introduced a new approach named AM-AFM tan δ imaging to characterize the mechanical dynamics of polymers at nanoscale.36 In the AM-AFM method, the cantilever is sinusoidally excited by an external force at an angular frequency, ω, near or at the fundamental value ω0. Under mechanical interaction with a sample, a part of the supplied power to the cantilever can be dissipated (denoted as Ploss), whereas a part of the supplied power is conserved during the interaction (denoted as Pstorage).28 These concepts in the AM-AFM, being similar to those in DMA, have led Proksch and Yablon to define the AM-AFM tan δ as36 tan δ =

Ploss ⟨F ·z⟩̇ ≡ − ts Pstorage ω⟨Fts·z⟩

(1)

where Fts is the tip/sample interaction force, z is the displacement of the cantilever from the equilibrium position, and ż is the cantilever velocity. A detailed description of the AM-AFM tan δ measurements can be found elsewhere.36,37,40 If ω is selected to be equal to ω0, eq 1 can be simplified to tan δ =

sin φ − A /A 0 cos φ

Figure 1. Loss tangent data for neat SBR and IR regions measured by the AM-AFM method in air and N2 environments in comparison with bulk data obtained by the DMA method.

(2) B

DOI: 10.1021/acs.macromol.8b01185 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

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RESULTS AND DISCUSSION The choice of SBR and IR here enabled the measurements of the viscoelastic dynamics of polymers in both glassy and rubbery regimes. We have previously reported that at the resonant frequency of the AM-AFM cantilever, ∼300 kHz, the glass−rubber transition temperature (Ttrans) of SBR was ∼30 °C above the Troom, whereas the Ttrans of IR was ∼25 °C below the Troom.37,44 This indicates that the viscoelastic responses of SBR and IR were in the glassy and rubbery regimes, respectively, at the Troom for the AM-AFM measurement. Such a difference is critical to distinguish the profile of tan δ along the interfacial region as illustrated in Figure 2. Figure 2a

Figure 3. Topography and phase shift images for (a, b) SBR/CB and (c, d) IR/CB samples measured by the AM-AFM method: 50 nm height scale and 30° angle scale for topography and phase shift images, respectively. Figure 2. (a) Schematic representing a gradient in the mobility slowing of polymer chains upon approaching the CB surface. (b) Profiles of tan δ on the neat polymer, interfacial, and CB regions predicted for glassy and rubbery responses of the neat matrix.

These images were taken in the net repulsive interaction at A/ A0 ∼ 0.5. It is clear that an IPR with a distinct value of phase angle and tan δ can be observed near the CB fillers for both rubber matrices, although it is difficult to recognize this IPR on the topography images. While it still remains challenging to connect the change in the phase angle to the quantity of polymer properties near the CB surface, the change in tan δ evidences that the viscoelastic dynamics of the IPR is different from the bulk behavior. It is noteworthy that such a change in the phase angle and tan δ observed near the nanofiller surface is generally not affected by the surface morphology.13,17 Panels d and h of Figure 4 show representative examples of the height and tan δ profiles taken through the same CB region of both topography and tan δ images for each sample, highlighted by solid lines on corresponding images. These profiles clearly indicate that a change in the tan δ near the CB region is not correlated to a change in the morphology of the sample surface. For example, there is a gradual increase in the height profile of the CB/IR topography image near the CB domain shown in Figure 4e, but the tan δ of the CB/IR interface shown in Figure 4g exhibits both the increased and decreased features, depending on the distance from the CB surface. Therefore, it is reasonable to conclude that the influence of the sample morphology on the tan δ imaging can be ignored; thus, the contrast in the tan δ through the IPR reflects the intrinsic dynamics of polymer chains near and far from the filler surface. This conclusion is indeed consistent with some previously reported results using AM-AFM phase imaging on various nanofiller/rubber interfaces.10,13,17 A more detailed analysis of the IPRs observed for CB/SBR and CB/IR tan δ images is presented in Figure 5. Figure 5a shows the histogram of the tan δ of a small area (shown in the inset) surrounding some CB fillers taken from Figure 4c. A smooth change of the tan δ value can be observed through this histogram. The tan δ is very small (∼0.1−0.2) on the CB

shows a schematic illustration of the segmental mobility of the IPR near a solid interface, such as CB, which is supposed to gradually slow down from the bulk behavior upon approaching the solid surface, as suggested in many reports.2,5,6,21 For a glassy matrix, tan δ of the IPR is expected to gradually decrease from the bulk value on the neat polymer region to a minimum value on the nanofiller (curve 1 in Figure 2b). In contrast, for a rubbery matrix, there might exist a glass−rubber transition point within the IPR, at which the viscoelastic response of the polymer changes from the rubbery behavior on the neat polymer region to the glassy behavior near the nanofiller interface. Similar to the temperature effect on the viscoelastic dynamics of polymers,37,39 the tan δ of the polymer is expected to reach a maximum value at the transition point before decreasing to a minimum value on the nanofiller (curve 2 in Figure 2b). We first performed the AM-AFM topography and phase measurements to characterize the dispersion of CB fillers within SBR and IR matrixes. Figure 3 reveals a good dispersion of the CB fillers in both rubber matrixes. The presence of individual nanofillers and small aggregations can be observed for both CB/rubber systems, where the size of the CB fillers was estimated to range from tens to hundreds of nanometers, in agreement with the size distribution of CB aggregates dispersed in various rubbers.45 Such a good dispersion of CB fillers within polymer matrices also indicates a favorable interaction between polymer chains and particle surface. To visualize the interfacial regions between nanofillers and rubbers, small areas surrounding some CB fillers were selected. The topography, φ, and tan δ measurements over such small areas of CB/SBR and SB/IR samples are provided in Figure 4. C

DOI: 10.1021/acs.macromol.8b01185 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

Figure 4. (a) Topography, (b) phase shift, and (c) tan δ images for CB/SBR interfacial regions and (d) surface profiles corresponding to the lines on the (a) topography and (c) tan δ images. (e) Topography, (f) phase shift, and (g) tan δ images for CB/IR interfacial regions and (h) height and tan δ profiles corresponding to the lines on the (e) topography and (g) tan δ images. The scan size for CB/SBR images was 300 × 300 nm2, and that for CB/IR images was 500 × 500 nm2.

Figure 5. Histograms of tan δ over small areas around (a) SBR/CB and (b) IR/CB interfaces: the insets show the corresponding tan δ images of these areas. The tan δ profiles are taken across the IPRs (corresponding to black lines on the inset images) for (c) SBR/CB and (d) IR/CB samples. Red dashed lines show the limits of the IPRs with the neat rubber and CB regions.

observation directly evidences that the segmental dynamics of polymer chains at the CB/SBR interface is slower than the bulk dynamics due to the interaction with the CB fillers. The slowing down is not sudden but gradual from the bulk dynamics upon approaching the filler surface. Therefore, it can be deduced that the segmental dynamics of the IPR in the CB/ SBR sample behaves like a gradient in the slowing down from the bulk dynamics to that of the CB domain, in line with some recent NMR and DRS results.6,15,16 Figure 5b shows the histogram of the tan δ of a small area (shown in the inset) surrounding some CB fillers in the IR matrix taken from Figure 4g. The smallest tan δ of ∼0.4−0.5 can be found on the CB domains. This value is again larger than the expected tan δ of the CB fillers, mainly attributed to the residual layer of IR on the CB surface, as in the case of the CB fillers in the SBR matrix. However, the tan δ of the IR near

region. This is because there is almost no viscoelastic interaction at the CB surface. In principle, the tan δ of the CB fillers should be nearly zero; however, as discussed in some previous reports,13,19 there always exists a portion of adsorbed polymers on the surface of nanofillers, which is difficult to be completely removed by cryo-sectioning. Consequently, the residual polymer might give rise to some contributions to the CB tan δ. The tan δ then gradually increases with increasing distance from the filler to the bulk region. An average tan δ value of ∼0.2−0.4 on the IPR can be estimated, which is relatively smaller than the peak value of ∼0.55 averaging over the neat SBR matrix. Because the viscoelastic response of the SBR exhibits the glassy behavior at the measured frequency of ∼300 kHz, a reduction in the tan δ of the CB/SBR interface with respect to that of the neat SBR region implies that the SBR/CB interface is more glassy than the neat SBR. This D

DOI: 10.1021/acs.macromol.8b01185 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules the CB fillers shows a much more complex behavior than that found at the SBR/CB interface. In fact, two interfacial layers near the CB fillers with distinct tan δ values in comparison with the peak value of ∼1.2 averaging over the neat IR region can be recognized. The first interfacial layer directly bound to the CB fillers exhibits a tan δ value from ∼0.6 to ∼1.0, relatively smaller than the bulk value, whereas the tan δ of the next interfacial layer becomes significantly larger than the bulk value, i.e., in a range of ∼1.5−2.5. Taking into consideration that the IR matrix exhibits a rubbery response at the measured frequency in the AM-AFM, the change in tan δ at the CB/IR interface can be explained as follows. The IR layer very close to the CB surface behaves as a glassy polymer having a relatively small tan δ value; further from this glassy layer, the IR dynamics becomes faster than that of the glassy layer and approaches the glass−rubber transition dynamics manifested by a substantial increase in the tan δ; after reaching the maximum value at the glassy−rubber transition point the tan δ decreases to the bulk value at the neat IR region. To the best of our knowledge, an interfacial polymer region showing such a distinct dynamic properties as observed here at the CB/IR interface represents the first experimental result in real space, which obviously confirms the existence of different dynamic layers in the IPR, although this dynamic behavior has been well described in various theoretical and simulation studies.5,21 Experimental profiles of the tan δ across both CB/rubber interfaces are plotted in panels c and d of Figure 5. We can see an excellent agreement between these experimental data and those predicted in Figure 2b. The tan δ data obtained here provided a direct evidence of the gradient in the slowing down of the segmental mobility of interfacial polymers upon approaching the CB interface, irrespective of the glassy or rubbery nature of the polymer matrix. In fact, the gradient behavior in the mobility slowing down of interfacial polymers manifested by an increase in Tg or relaxation time has been previously confirmed in NMR and DRS experiments.5,6,15,16 However, unlike the previous experiments where some models were needed to describe the gradient behavior, the AM-AFM tan δ imaging could directly visualize the mobility gradient of the interfacial polymers. In addition, from the profiles in panels c and d of Figure 5 we could directly calculate the length scale of the IPRs. We suppose that the limits of these regions are at the points where tan δ starts deviating from the values averaging over the plateau regions of the CB and neat rubbers, as highlighted with the dashed lines in Figures 5c and 5d. The length scale of the CB/SBR interface was found to be ∼24 nm, whereas the length scale of the CB/IR interface was ∼32 nm. These length scales are about 3−4-fold larger than the probe radius; thus, it is unlikely that these values were convoluted with the probe size. The difference in the length scale of the IPR might be due to a difference in the interaction strength between CB particles and two rubbers. Many reported data have indeed revealed that increasing the interaction between the nanofiller surface and polymer matrix can lead to an increase in the length scale of the interfacial region.1,4,18 The length scale can be also affected by the state of the polymer matrix, i.e., glassy or rubbery, at the measured temperature and frequency.5 Although this length scale is relatively larger than that of ∼5−10 nm generally reported using DRS and thermal analysis techniques,15,16,26 it is in good agreement with that observed on different filler/rubber systems using NMR and AFM methods.5,12,13,17,46

It is also interesting to note that in the case of CB/IR we could observe the existence of two interfacial regions, in which the glassy layer was ∼14 nm, in excellent agreement with the result obtained by Berriot and co-workers using the NMR method in considering that the measured temperature was ∼25 °C above the Ttrans.5,46 To explain such a large glassy layer, Long and co-workers have proposed a theoretical model, also known as percolation model, which was developed on the basis of the heterogeneous nature of the dynamics in glass-forming materials.47,48 In this model, the thickness of the glassy layer (hglass) was proposed to be a function of the difference between the measured temperature (Tmeas) and the Ttrans hglass

ij yz Ttrans zz = ξjjj j Tmeas − Ttrans zz k {

ν

(3)

where ξ is the length of the order of the nanometer that depends on the polymer/filler interaction and ν = 0.88 is the exponent for the correlation length. For the CB/IR sample, Ttrans = 273 K, Tmeas = 298 K, and hglass = 14 nm; we can therefore estimate that ξ = 1.7 nm, which is in line with the predicted value in the percolation model.5,47,48

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CONCLUSIONS We have directly mapped the viscoelastic dynamics of CB/SBR and CB/IR interfaces by taking advantage of the AM-AFM tan δ imaging at nanoscale. In both cases, the tan δ images have shown the gradient nature in the slowing behavior of the polymer dynamics close to the nanofiller interface. The length scales of these IPRs have also been determined to be ∼24 and ∼32 nm for CB/SBR and CB/IR interfaces, respectively. Our study can therefore help to answer the long-existing problem of visualizing the interfacial dynamics in polymer nanocomposites in real space.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (K.N.). 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 We thank Dr. R. Proksch in Asylum Research, Oxford Instruments, for fruitful discussions during AM-AFM loss tangent measurements and Dr. T. Igarashi at Bridgestone Corp. for preparing the samples. This work was supported by JSPS Grant-in-Aid for Scientific Research (B) (26288097).

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REFERENCES

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