Interphase Percolation Mechanism Underlying ... - ACS Publications

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Cite This: J. Phys. Chem. C XXXX, XXX, XXX−XXX

Interphase Percolation Mechanism Underlying Elastomer Reinforcement Siwu Wu, Min Qiu, Zhenghai Tang, and Baochun Guo* Department of Polymer Materials and Engineering, South China University of Technology, Guangzhou 510640, P. R. China S Supporting Information *

ABSTRACT: Glassy interphase has been claimed to be of vital importance for mechanical reinforcement of elastomer nanocomposites (ENCs), but the evolution of interphase topology in correlation to reinforcement percolation remains uncertain. Here, an accurate interfacial regulation strategy upon implementing an interphase percolation mechanism is exploited to realize percolation of mechanical performance toward striking elastomer reinforcement. Architecture design of interfacial metal−ligand bridges accomplishes firm anchoring between elastomer skeleton and carbon nanodots, leading to the formation of interfacial metal-enriched regions. The volume fraction of the interfacial region systemically enlarges upon increase of interfacial bridges, which finally overlaps with neighboring domains to form a penetrating interphase. The topological evolution of the interfacial region is quantitatively monitored upon small-angle X-ray scattering and dielectric measurements, which exhibits a similar percolation behavior in sync with that of macroscopic mechanical performance. Furthermore, the interphase exhibits much slower relaxation dynamics than in bulk polymer, which significantly improves the network rigidity and hence accounts for the prominent elastomer reinforcement. This investigation corroborates that the formation of penetrating interphase may be an executable mechanism to induce the reinforcement percolation of ENCs. We further envision that the implementation of interphase percolation mechanism can be a universal avenue to afford rationalized optimization of ENCs.

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INTRODUCTION Elastomer nanocomposites (ENCs) are strategically important and indispensable in a great variety of industrial and high-tech applications. In particular, nanoparticles reinforcement is a generic and well-accepted solution toward high-efficiency strengthening of elastomers, which has been validated by extensive studies and already implemented in practical applications.1,2 The exploration on the reinforcement mechanism of ENCs has long been a fundamental and essential theme in the community.3,4 To date, extensive experimental and theoretical efforts have been dedicated to specify the reinforcement mechanism of ENCs in relation to filler−filler and matrix−filler interactions beyond hydrodynamic effects.5,6 Generally, in the case of ENCs with sufficiently high loading of nanoparticles, the formation of percolating filler network and an interfacial layer with glassy or much restricted dynamics in the vicinity of nanoparticles play a crucial role in elastomer reinforcement at small strain.7,8 As for the large strain regime, the percolating filler network can enable sufficient chain alignment and orientation between two adjacent nanoparticles to significantly improve the resistance to external stress, leading to highly efficient strengthening for ENCs at large strain.9,10 Besides, several models have been proposed based on computer simulations, which indicated the formation of penetrating interphase upon overlap of discrete interfacial regions may account for the reinforcement percolation of ENCs.11−13 This © XXXX American Chemical Society

assumption has also been adopted to explicate the sol−gel transition behavior in polymer melts and liquid polymer nanocomposites.14−16 Nevertheless, in most of ENCs system, the interfacial layer bridging the interpenetrating filler network is superficially associated with mechanical reinforcement by facilitating the propagation of external stress.17,18 A distinct experimental correlation between microscopic penetrating interphase and macroscopic mechanical performance of ENCs is fundamentally important but still lacking. In ENCs system with high filler loading, the substantial contribution of penetrating interphase to macroscopic reinforcement and microscopic chain dynamics may easily be obscured from the significant effects of the filler network.19,20 As for the systems with moderate filler loading, it is very difficult to precisely regulate the fraction of interfacial region by in situ interfacial modification, although it may be realized through some complex and time-consuming pretreatment procedures.14,16,21 Metal−ligand coordination interactions are prevalent in biomimetic systems by the merits of precise bond directionality, well-defined stoichiometry, and a broad range of molecular parameters (different metal ions, counterions, and ligands), endowing systems with a number of desirable material Received: November 14, 2017 Revised: December 3, 2017 Published: December 6, 2017 A

DOI: 10.1021/acs.jpcc.7b11239 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C

solution. The transparent solution was subsequently placed into a household microwave oven (700 W) and heated for 3 min. After cooling to room temperature, the obtained orange-red solid was diluted with few distilled water and then washed with excess anhydrous ethanol repeatedly for several times to remove byproducts and unreacted reagents. The purified CDs were dried in a vacuum at 50 °C for further characterization and application. Preparation of ENR Samples. Desired amounts of CDs and FeCl3 were dissolved in 5 mL of acetone. After adequately stirring, the mixture was dropwise admixed with ENR in an internal mixer (50 rpm) at 80 °C for 15 min. After cooling to room temperature, the compounds were incorporated with curing ingredients on an open two-roll mill and then subjected to compression at 143 °C for the optimum curing time, which was determined by a vulcameter. The specific formulation is listed as ENR 100 g, zinc oxide 5 g, stearic acid 2 g, Ncyclohexyl-2-benzothiazole sulfonamide (CZ) 1.5 g, sulfur 1.5 g, CDs 5 g, and FeCl3 variable. In the context, the sample code of EC5Fx refers to ENR sample contained 5 phr (parts per hundreds of gum) of CDs and x phr of FeCl3. For comparison, the counterparts excluding CDs (referred to as EC0Fy, where y denotes y phr FeCl3) were also prepared according to the aforementioned process. Characterization. The morphology and size of CDs were captured using a Philips FEI-Tecnai G2 F30 S-Twin TEM microscope at an accelerating voltage of 300 kV. The sample for HRTEM observation was prepared by dropping CDs solution onto a copper grid coated with carbon film. Fourier transform infrared spectroscopy (FTIR) spectra were recorded using a Bruker Vertex 33 FTIR spectrometer. UV−visible (UV−vis) absorption spectra were recorded on a PERSEE TU1810DSPC UV−vis spectrophotometer. Fluorescence (FL) emission spectra were measured on an Edinburgh FL920 fluorescence spectrometer. The cross-linking densities of the samples were determined by the equilibrium swelling method and was calculated using the Flory−Rehner equation.30 All samples were immersed in tetrahydrofuran until equilibrium swelling (72 h). Tensile stress−strain behavior was measured using a Gotech AI-7000 S servo control system universal testing apparatus at 25 °C. For uniaxial tensile tests, the strain rate was 500 mm/ min, following ISO standard 37-2005. For cyclic tensile tests, all measurements were carried out at 25 °C with a strain rate of 100 mm/min, and the stress−strain hysteresis was calculated from the area between the loading and unloading loop. X-ray scattering experiments were performed on Bruke AXS Nanostar grazing incidence small-angle X-ray scattering (SAXS) apparatus using Ga Kα X-rays generated from a METALJET Xray source generator operated at 40 kV and 40 mA. The scattering data of the ENR samples were fit to

properties, including self-healing, dynamic reversibility, high toughness and hardness, and mechanical tunability.22,23 Hence, metal−ligand coordination interactions are uniquely suitable to engineer robust and tunable mechanical performance for ligand-containing polymers by simply selecting the type and amount of added metal cross-linkers.24−26 Previously, we have successfully constructed interfacial metal−ligand bridges between a commercial rubber with pendent pyridine moieties and the surface catechol groups on reduced graphene, resulting in a robust elastomer with integrated improvement of strength, toughness, and stretchability.26 However, the influence of interfacial metal−ligand bridges on the macroscopic mechanical performance of ENCs remains as qualitative. In this work, our focus is dedicated to disclose the microscopic interfacial evaluation mechanism underlying the elastomer reinforcement by means of experimental trial, aiming at providing new insights into the vital role of penetrating interphase for elastomer reinforcement. For this purpose, an accurate interfacial regulation strategy is exploited to realize percolation of mechanical performance toward striking elastomer reinforcement. Specifically, epoxidized natural rubber (ENR), a commercial ligand-containing elastomer, is employed as an elastomer matrix in order to facilitate tunable interfacial regulation by virtue of the abundant oxygenic moieties in the chain skeleton. Meanwhile, as a newly emerging class of nanoparticles, carbon nanodots (CDs) possess abundant surface functional moieties with lone electron pairs, such as amino, amide, and hydroxyl groups, which impart them with unique capability of interacting with metal ions.27,28 Accordingly, architecture design of interfacial metal−ligand bridges is implemented to directionally establish firm anchoring between oxygenic moieties in the skeleton of ENR and surface functional moieties of CDs, leading to the formation of interfacial metal-enriched regions. These metal-enriched domains can serve as physical cross-linkers to improve the network rigidity in bulk and are crucial for the prominent mechanical reinforcement of the system. Subsequently, the volume fraction of interfacial region systemically enlarges upon increase of metal−ligand bridges, which finally overlaps with neighboring domains to form a penetrating interphase with distinctive relaxation dynamics. Through an integrated approach, we have therefore corroborated the volumetric enlargement of interfacial region exhibited a similar percolation behavior in sync with that of macroscopic mechanical performance. We envision that a systematic study will contribute to gain deeper insights into the vital role of penetrating interphase for elastomer reinforcement and potentially afford rationalized optimization of ENCs.

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EXPERIMENTAL SECTION Materials. Epoxidized natural rubber (ENR, epoxidation degree = 50%) was manufactured by Agricultural Products Processing Research Institute, Chinese Academy of Tropical Agricultural Science, Zhanjiang, China. Critic acid (CA, ≥99.5%), 1,2-ethylenediamine (EDA, 99%), anhydrous iron(III) chloride (FeCl3, 98%), and propylene oxide (PO, 99%) were purchased from Beijing InnoChem Science & Technology Co., Ltd., Beijing, China. All rubber additives were industrial grade and used as received. Synthesis of Amine-Passivated Carbon Nanodots. CDs were synthesized by microwave-assisted pyrolysis process according to previously reported method.29 Typically, CA (1.0 g) and EDA (0.313 g) were dissolved in 10 mL of phosphate

I(q) = IKT(q) + C

where IKT(q) is the Kinning−Thomas modified hard-sphere model (see Supporting Information for details) and C is a constant accounting for the background scattering.31 Broadband dielectric relaxation spectroscopy (BDRS) measurements were conducted using a Novocontrol ALPHAANB broadband dielectric/impedance spectrometer equipped with a Quatro Cryosystem temperature controller. Film samples with a thickness of ∼0.5 mm were sandwiched between parallel gold-plated electrodes with a diameter of 25 mm. Measurements swept over a frequency were recorded over B


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Figure 1. (a) HRTEM image of CDs. Inset: size distribution histogram of CDs. (b) FTIR spectra of citric acid and CDs. (c) UV−vis absorption and FL emission spectra of CDs solution.

behavior has been confirmed within the λex range from 300 to 420 nm, which implies that both size and surface state of the asprepared CDs are uniform.38,40 Interaction between Amine-Passivated Carbon Nanodots and Ferric Ion. As a typical d-block ion, ferric ion is liable to open pathways through electron transfer involving the metal center. It is not surprising, therefore, that abundant surface functional groups with lone electron pairs, such as amino, amide and hydroxyl groups, of the as-prepared CDs are likely to possess excellent and firm binding affinity toward ferric ions. UV−vis spectra and FL quenching assay were employed to uncover the interaction between CDs and ferric ions. Upon titration of CDs with ferric ions, the π−π* transition band exhibits a red-shift to 246 nm while the n−π* absorption peak appears to shift to lower wavenumber of 344 nm (Figure 2a).

a frequency (f) ranged from 1 to 107 Hz at isothermal conditions that varied from 10 to 100 °C in 10 °C steps. The obtained dielectric loss (ε″, imaginary part of dielectric permittivity) was analyzed by the classic Havriliak−Negami (HN) function with a conductivity term:32−34 ⎛ σdc ⎞ N * ε (ω) = −i⎜ ⎟ + ⎝ ε0ω ⎠

n

⎡

k=1

⎣

∑ ⎢⎢ε∞k +

⎤ ⎥ (1 + (iωτHNk)αk )βk ⎥⎦ Δεk

where σdc is the dc conductivity, ε0 is the vacuum permittivity, N is an exponential factor (0 < N ≤ 1), ε∞k is the relaxed value of the dielectric constant of kth process, Δεk is the dielectric strength of the kth process, ω is the angular frequency, and τHNk is the HN characteristic relaxation time of the kth process. αk and βk are the shape parameters of the kth process. The relaxation time of maximum loss is referred to as τmax, which can be determined from the HN fitting parameters:

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−1/ α 1/ α ⎛ ⎛ αβπ ⎞ απ ⎞ τmax = τHN⎜sin ⎟ ⎜sin ⎟ ⎝ 2 + 2β ⎠ ⎝ 2 + 2β ⎠

RESULTS AND DISCUSSION Characterization of Amine-Passivated Carbon Nanodots. Quasi-spherical, uniformly distributed dark features, corresponding to the synthesized CDs, are observed in the HRTEM image (Figure 1a). The average diameter of CDs is determined as 2.2 ± 0.5 nm, which is coincident with our previous studies.35,36 Surface functionality of CDs was investigated using FTIR measurement, providing insights into the chemical structure. As shown in the FTIR spectrum of CDs (Figure 1b), the sharp absorptions centered at 1655 and 1560 cm−1 are typical peaks for the CO stretching and N−H blending vibration of the amide group.37 Also, the absorption centered at 1190 cm−1 and the broad absorption band at 3150− 3340 and 3100−3700 cm−1 are assigned to stretching vibrations of the C−NH bond and N−H and O−H stretching vibration of secondary amine and hydroxyl groups, respectively.28,38 These peaks clearly indicate that hydroxyl, amide, and secondary amine groups are the main passivated moieties on the surface structure of CDs. Two distinct absorption peaks centered at 238 and 350 nm can be observed in the UV−vis spectrum of CDs aqueous solution, which are assigned to the π−π* transition of the CC aromatic structure in the carbogenic core and the n−π* transition of the carbonyl groups, respectively.39 Under UV irradiation (365 nm), the CDs solution emits bright blue fluorescence with a maximum emission peaks centered at round 455 nm (Figure 1c). Notably, excitation-independent FL

Figure 2. (a) UV−vis absorption spectra of CDs, FeCl3, and their mixture at various CDs:FeCl3 mass ratio. (b) Evolution of FL emission spectra of CDs upon titration of various amounts of FeCl3. Inset: the linear relationship between F0/F − 1 and the concentration of Fe3+, where F0 and F are FL intensity at 455 nm in the absence and presence of Fe3+, respectively.

The former one may be assigned to a delocalization of π electrons induced by the positively charged Fe3+ center, and the latter one most likely suggests a ligand-to-metal charge transfer between Fe3+ center and the surface functional groups with lone pair electrons on CDs.41,42 Moreover, it is widely known that the coordinative interactions of CDs with Fe3+ are supposed to activate a sensitive fluorescence quenching of CDs. As shown in Figure 2b, the maximum FL intensity at 455 nm of CDs solution monotonically quenches upon gradual titration of Fe3+. The quenching efficiency (F0/F − 1) displays a good linear relationship versus the concentration of Fe3+ in the range 0−80 μM. These results are in accordance with previous reports, which confirms that the formation of coordination bonds between Fe3+ and surface amino, amide, and hydroxyl groups of CDs can bring into a nonradiative C


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Figure 3. (a) Representative stress−strain curves of (a) EC5F1.5 and control samples and (b) EC5Fx samples. (c) Tensile loading−unloading curves of EC5Fx samples. (d) Dependence of hysteresis area and tensile modulus (stress at 100% strain) of EC5Fx (solid symbol) and EC0Fx (open symbol) samples on the FeCl3 content.

are essential for pronounced reinforcement in ionomers.46 Even so, the total cross-linking density of EC0F1.5 increases upon formation of Fe3+−epoxy coordinations (as shown in Figure S2), which inevitably leads to a significant decline in stretchability. In addition, the as-prepared CDs possess abundant active-hydrogen surface moieties as aforementioned, which may also generate hydrogen-bonding affinity with the oxygenic moieties in the elastomer skeleton and contribute to the mechanical reinforcement of ENR matrix. To explicitly address this issue, we employ propylene oxide (PO) as the model compound, which is the main constitutional unit of ENR proposed to interact with CDs surface groups. As indicated by UV−vis spectra (Figure S3a), this interfacial hydrogen-bonding affinity is too weak to be detected. Hence, the reinforcement effect of single incorporation of CDs on ENR matrix is negligible as shown in Figure S3b (detailed mechanical properties are tabulated in Table S1). Compared with the parental formulations, EC5F1.5 sample enables remarkable reinforcement (∼7-fold increase) of tensile modulus upon construction of interfacial metal−ligand bridges (the corresponding mechanical data are tabulated in Table S1). We believe the Fe3+ ions could serve as intermediaries to anchor matrix chains in the vicinity of particulate CDs upon formation of interfacial metal−ligand bridges. Consequently, interfacial regions with restricted chain mobility are formed, which may impede the network relaxation and significantly increase the network rigidity in bulk.12 The dissociative character of ligand exchange in bulk samples makes the CDs−Fe3+−epoxy interfacial bridges similar to covalently bound interfaces which generally lead to declined stretchability and enhanced modulus.22,47 By virtue of this accurate interfacial regulation

electron transfer involving partial transfer of an electron in the excited state to the free orbital of Fe3+.27,28 Hence, the radiative electron transfer of FL emission is bound to be disrupted and restrained, leading to obvious FL quenching. In addition, we have confirmed that strong metal−oxygen coordinations were enabled upon incorporation of metal ions with ENR matrix in our previous research.24 Hence, it is reasonably believed that Fe3+ ions can serve as intermediaries to directionally anchor ENR chains in the vicinity of particulate CDs upon formation of interfacial metal−ligand bridges. Contribution of Interfacial Metal−Ligand Bridges to Mechanical Properties. It is now well-recognized that the alteration of matrix−particle interactions gives rise to a wide tunability in macroscopic properties for polymer nanocomposites.43 To disclose the contribution of interfacial metal-enriched bridges to the mechanical performance of elastomer, investigation was started by taking EC5Fx samples as a mode system to gain an insight into the crucial role of interfacial regions. Typical tensile stress−strain curves of ENR composites are depicted in Figure 3a. Obviously, the direct implementation of metal−ligand bonds into the ENR skeleton can hardly bring an improvement on the integral mechanical performance but impair the stretchability apparently (Figure S1 and Table S1). This seems contrary to previous studies on ionic elastomers, in which significant improvement on mechanical performance has been achieved through the incorporation of chelate metallic ions.44,45 However, the dosages of Fe3+ in the EC0Fy system are too low to reach the stoichiometric ligand/ metal molar ratio (epoxy:Fe3+= 6:1);24 the electrostatic attractions are therefore unable to induce efficient aggregation of ion pairs to form large multiplets or even ionic clusters which D


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Scheme 1. (a) In the EC5Fx System, Architecture Design of Interfacial Metal−Ligand Bridges Implements Firm Anchoring between Elastomer Skeleton and Carbon Nanodots, Leading to the Formation of Interfacial Metal-Enriched Region; (b) Subsequently, the Volume Fraction of Interfacial Region Systemically Enlarges upon Increase of Metal−Ligand Bridges; (c) When Fe3+ Concentration Exceeds Percolation Threshold, the Discrete Metal-Enriched Domains Grow to Overlap with Neighbors

Figure 4. Temperature dependence of loss tangent (tan δ) of (a) EC5F1.5 and control samples and (b) EC5Fx samples. (c) X-ray scattering intensity of EC5Fx samples vs scattering vector q. The dashed line is the best fit of the KT modified hard-sphere model to the scattering data. (d) Dependence of Rpc, Rmd, and Φi for EC5Fx samples on the content of FeCl3.

interactions (Figure S4, detailed data are listed in Table S2) indicates that the present strategy exhibits superior reinforcement efficiency for elastomers in a relatively low filler content. Interestingly, with a further survey of the mechanical behaviors of EC5Fx samples, a sharp transition in bulk mechanical performance at ∼1.25 phr FeCl3 comes into sight (Figure 3d). At this point, the sample transits from a

strategy, it is possible to modulate the overall mechanical performance of ENR samples by simply adjusting the amount of metal concentration to rationalize the structural topology of metal−ligand interfacial region. As expected, the tensile modulus and hysteresis of EC5Fx samples increase monotonously with Fe3+ content (Figure 3b,c). Careful survey of other reported ENR-based composites with various interfacial E


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Figure 5. (a) Temperature and frequency dependence of dielectric loss ε″ for EC5F1.5. Deconvolution results for the dielectric loss ε″ of EC5F1.5 at (b) 20 °C and (c) 70 °C. Solid, dashed, and dash-dotted lines represent the total fitting curve (HN function and conductivity), the individual processes, and the conductivity contribution, respectively. (d) Temperature dependence of average relaxation times for segmental and interfacial process of EC5F1.5 and control samples. Dashed lines are VFT fits to the data. The parameters for these fits are listed in Table S5. (e) Frequency dependence of dielectric loss ε″ for EC5Fx samples at 70 °C. (f) Volume fraction of interphase defined as ΔεInterphase/ΔεTotal polymer for EC5Fx samples and the relative dielectric strength of segmental process Δεbulk for EC5Fx (solid symbol) and EC0Fx (open symbol) samples as a function of FeCl3 content. All Δε values are obtained from the dielectric loss spectra at 70 °C.

system is fundamentally in accordance with previous results based on computer simulation, in which the reinforcement percolation of polymer nanocomposites was attributed to the formation of penetrating interphase upon overlap of discreate interfacial regions.11−13 Exploration of Interphase Percolation Behavior. It is no doubt that the formation and population of interfacial metal−ligand bridges in the EC5Fx system not only evoke a sharp transition behavior of macroscopic mechanical performance but also generate a distinctive network topology compared with EC0Fy counterparts. In order to better understand the origin of reinforcement percolation of elastomer matrix at molecular scale, the evolution of interphase topology and relaxation dynamics upon increasing interfacial metal−ligand bridges should be specialized. We started our investigation from probing the response of overall network relaxation behavior upon dynamic mechanical loading. As shown in Figure 4a and Figure S5, the temperaturedependent loss tangent (tan δ) curves of EC0Fy feature a gradual decline of peak value, along with an increase of glass transition temperature (Tg). This indicates that the direct incorporation of Fe3+ ions with the ENR matrix can only generate partial and discrete coordination complexes which impair the mobility of bulk network. As for the EC5Fx system, the variation trend of tan δ curves is similar to the EC0Fy counterparts at lower FeCl3 dosage. However, as FeCl3 content exceeds percolation threshold (∼1.25 phr), another legible relaxation process at higher temperature (∼55 °C) is observed for EC5Fx samples (Figure 4b). This is reasonable because the

comparably stretchable elastomer into a relatively plastic leather, leading to a striking improvement of modulus and hysteresis. A similar phenomenon has been observed in some reported ligand−polymer based systems, stemming from the drastic reduction of uncoordinated ligand moieties as the ligand/metal ratio approached equilibrium which slowed down the chain relaxation and significantly increased the network rigidity.22,47 As aforementioned, the content of CDs in the EC5Fx system is constant, and the dosage of Fe3+ in the current study is far from affording equilibrium metal−ligand network structural topology. Hence, we propose a hypothesis to interpret the sharp transition behavior of bulk mechanical performance in terms of the interphase percolation mechanism. At low FeCl3 content (less than 1.25 phr), only a small number of polymer chains are anchored on CDs to form thin interfacial metal-enriched layers with restricted mobility (Scheme 1a). These discrete domains could only provide limited reinforcement for the ENR matrix. Upon increase of Fe3+ concentration, the volume fraction of interfacial regions gradually enlarges (Scheme 1b), leading to a systemic response of mechanical performance of the system. When FeCl3 content exceeds percolation threshold (∼1.25 phr), these metal-enriched domains grow to overlap with neighbors, giving rise to a penetrating interphase with restricted mobility, which is separated from the bulk polymer (Scheme 1c). This interphase, associated by metal-enriched domains which serve as solid-like cross-linkers, significantly improves the network rigidity significantly and is crucial to the reinforcement percolation of EC5Fx samples.44,48 This mechanism involved in current F


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are depicted in Figure 5a. As the testing temperature approaches to the bulk glass transition, the segmental relaxation process associated with the cooperative motion of bulk chain segments appears and subsequently shifts to higher frequencies upon temperature increase due to thermal activation. This is also observed in the spectra of controlled samples (EC5F0 and EC0F1.5, as shown in Figure S7), and only one relaxation process is confirmed within the current temperature window. In contrast, a pronounced shoulder process with much slow relaxation emerges at higher temperature in the spectra of EC5F1.5. A similar phenomenon has been observed in previously reported polymer−nanofiller systems with various kinds of interfacial interactions, which assigned this slow process to the relaxation of interfacial chains with restricted mobility.53,54 However, the attribution of the observed slow process is still controversial because some other researches have referred it to the evolvement of Maxwell−Wagner−Sillars (MWS) or interfacial polarization due to the gathering of charges at the internal phase boundaries.55,56 Although MWS polarization exists in all heterogeneous systems, the intensity of this low-frequency process usually decreases upon increase of temperature.57 Apparently, the low-frequency shoulder of the EC5F1.5 sample is a typical temperature-dependent process with tenable peak intensity; hence, it is reasonable to assign this process (thereafter referred to as the interfacial process) to the relaxation of interfacial metal-enriched region. In order to distinguish the contribution of each relaxation process to the complex dynamics of EC5F1.5 sample, the phenomenological Havriliak−Negami (HN) function is introduced to analyze the corresponding dielectric loss spectra. At low temperature, all spectra with only one detectable relaxation process can be well described with single HN function plus a conductivity contribution. Figure 5b shows a typical fit of the dielectric data for EC5F1.5 at 30 °C to single HN function. However, with the emergence of interfacial process at higher temperature, the dielectric loss spectra of EC5F1.5 samples have to be deconvoluted with two HN functions plus a conductivity term, as depicted in Figure 5c. The obtained relaxation time τmax of segmental process for ENR samples are plotted as a function of reciprocal temperature (Figure 5d), which reveals a Vogel−Fulcher−Tammann (VFT) dependence as

enlargement of interfacial metal-enriched domains is bound to overlap with neighbors, leading to the formation of penetrating interphase with restricted relaxation behavior. The striking similarities in the dependence of both static and dynamic mechanical performance on the Fe3+ content suggest that the reinforcement percolation behavior for the EC5Fx system should originate from the formation of penetrating interphase, as aforementioned. To further provide a quantitative evaluation of interfacial geometry, SAXS measurements were performed to monitor the topological evolution of interfacial metal-enriched domains. Representative X-ray scattering profiles for ENR composites are displayed in Figure 4c. In the case of the EC5Fx system, the intensity of X-ray scattering as a function of scattering vector shows a distinct peak at q ∼ 1.5 nm−1, which should be attributed to the feature of interfacial metal-enriched domains. As for the EC0Fy system, no visible peak can be distinguished in the testing range (Figure S6). The peak intensity of metalenriched domains significantly enhances upon increase of Fe3+ content due to the incremental electron density contrast.31 The scattering data for the EC5Fx system were fitted with the KTmodified model to facilitate the analysis of scattering peak.49,50 Four independent fitting parameters are induced to characterize metal-enriched domains: the radius of particulate center Rpc, the minimum radius of closest approach between two adjacent metal-enriched domain centers Rmd, the average volume of domains Vad, and the peak amplitude of the scattering maxima A. Obviously, this model provides a good fit to the scattering data, and the corresponding parameters are tabulated in Table S3. As shown in Figure 4d, the increase of Fe3+ content performs a declining effect on the size of particulate center Rpc but an increase in Rmd, simultaneously. These results are plausible as increasing Fe3+ content is certain to not only facilitate the dispersion of CDs particles, leading to disintegration of CDs aggregates and a decline of Rpc, but also enlarge the thickness of impenetrable metal-enriched shell surrounding CDs particles, resulting in an increase of Rmd. The calculated Rpc value of the EC5F2 sample is very close to the size of CDs as indicated by TEM image (Figure 1a), which implies a uniform dispersion state of CDs is achieved. Moreover, as the thickness of metal-enriched shell is on the order of Rmd − Rpc according to the pervious study, the volume fraction of interphase (Φi) can be determined as

⎛ B ⎞ τmax = τ0 exp⎜ ⎟ ⎝ T − T0 ⎠

4 Φi = π (R md 3 − R pc 3)NpcVem−1 3

where τ0 is the infinite relaxation time, B is an empirical parameter and inverse to the apparent activation energy,50 and T0 is the Vogel temperature. As shown in Figure 5d, the segmental relaxation dynamics of EC0F1.5 is obviously slower than that of EC5F0, demonstrating a restricted chain mobility of bulk network is achieved upon direct implementation of metal−ligand bonds into the ENR skeleton. In addition to the pronounced restriction of segmental relaxation dynamics, the architecture design of interfacial metal−ligand bridges in EC5F1.5 sample contributes to an interfacial process with orders of magnitude slower dynamics than that for bulk chains. Coincidentally, the dissociation of reversible complexes, such as multiple hydrogen-bonding and metal−ligand combinations, can also lead to a high-temperature dynamical process. Such process is a typical chemical relaxation for which the temperature dependence of relaxation time generally exhibits Arrhenius behavior.58,59 Meanwhile, its relaxation strength

where Npc is the number of CDs (the molar mass of CDs is assumed as 2000, referring to the previous research51) and Vem is the volume of ENR matrix. For the EC5Fx system, a transitive enlargement of Φi is found as the FeCl3 content exceeds 1.25 phr, performing in a typical percolation behavior which provides an explicit evidence to support our aforementioned hypothesis. So far, we have provided convincing evidence to demonstrate that the reinforcement percolation of the EC5Fx system stemmed from the formation of penetrating metal-enriched interphase. Nevertheless, the dynamics of the EC5Fx system are still necessary to provide a holistic understanding of the macroscopic percolation behavior of these materials. From this point of view, BDRS measurements were conducted due to the high accuracy at probing and decoupling complex dynamics of polymer composites.52 The temperature-dependent dielectric loss spectra for EC5F1.5 sample over a wide range of frequency G


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ΔεInterphase/ΔεTotal polymer (ΔεTotal polymer is determined as ΔεInterphase + ΔεBulk) can be used as a quantitative indicator of the volume fraction of interfacial region.52,53 The corresponding values of ΔεInterphase/ΔεTotal polymer are tabulated in Table S6, which are very close to those determined by SXAS measurements. As expected, the volumetric enlargement of metalenriched interphase for the EC5Fx system performs a percolation behavior with a threshold FeCl3 content around 1.25 phr. This further confirms our aforementioned hypothesis that the reinforcement percolation of the EC5Fx system originates from the formation of penetrating metal-enriched interphase. The architectural design of interfacial metal−ligand bridges into EC5Fx samples is inclined to anchor matrix chains in the vicinity of particulate CDs, leading to the formation of interfacial metal-enriched region with much slow relaxation. Upon increase of Fe3+ concentration, the volume of interfacial region grows to overlap with neighboring domains, giving rise to a penetrating interphase with distinctive dynamics, which significantly improves the network rigidity and crucial to the reinforcement percolation of the EC5Fx system.

correlated to the number of association complexes decreases significantly with increasing temperature.60,61 However, the analysis of the temperature dependence shows that the hightemperature relaxation of EC5F1.5 sample follows a VFT dependency (Figure 5d). Moreover, only a slight decrease of the dielectric strength of this relaxation at elevated temperature is characteristic (detailed values have been tabulated in Table S4). Hence, it is reasonable to assign the high-temperature process to the dynamic glass transition of the interfacial chains in the present system. By extrapolating VFT fit to the temperature at which τ = 100 s, an estimated Tg100s for corresponding relaxation process can be obtained. The Tg100s of segmental and interfacial relaxation for EC5F1.5 sample are determined as ∼10 and ∼49 °C, which are very close to those evaluated by DMA measurement (Table S5) and hence implies the metal-enriched interphase with restricted mobility possesses a “glassy” behavior compared to bulk polymer. To further confirm the improved network rigidity of ECF5F1.5 resulted from the formation of penetrating metal-enriched interphase, the apparent activated energies (Ea = BR/([1 − T0]/Tg100s)2) of all relaxation processes are calculated and tabulated in Table S5. Apparently, the Ea of the segmental relaxation for EC0F1.5 sample is much higher than that of EC5F0, indicating the direct implementation of metal−ligand bonds into the ENR skeleton can give rise to a significant restriction of bulk network relaxation. For the EC5F1.5 sample, the formation of penetrating metal-enriched interphase not only leads to a further increase of the Ea of segmental relaxation, the relaxation of penetrating interphase also possesses the highest Ea as expected. This provides convincing evidence that the formation of penetrating metal-enriched interphase with restricted mobility improves the overall network rigidity. Besides, ascertaining the influence of interphase content on the overall relaxation behavior of the EC5Fx system is critical to the understanding the origin of macroscopic mechanical percolation. In principle, the confinement effect of interfacial metal−ligand bridges is expected to be increasingly important upon increase of Fe3+ concentration. The isothermal dielectric loss spectra of the EC5Fx system are depicted in Figure 5e. A significant broadening and shifting to lower frequencies of segmental process are observed with increasing Fe3+ concentration. The gradual broadening suggests developmental heterogeneous dynamics, and the slowing down is due to increscent steric hindrance,52 both of which qualitatively indicate the increasing restrictions of bulk relaxation stemmed from the enlargement of Fe3+-enriched domains upon increase of Fe3+ content. In addition, the dielectric strength (Δε) extracted from HN fits is proportional to macroscopic polarization dominated by the number density of dipole moment for a given process. Therefore, comparison of dielectric strengths associated with different relaxation process can quantify the variation of both bulk polymer (ΔεBulk) and interfacial region (ΔεInterphase).52 Upon direct implementation metal−ligand coordination into polymer backbone, ΔεBulk of the EC0Fy system is proportional to Fe3+ content due to the gradual increase of the overall dipole moment (Figure 5f). Notably, several interesting features can be found regarding the influence of Fe3+ content on the Δε of relaxation processes for EC5Fx samples. For segmental relaxation, ΔεBulk of the EC5Fx system exhibits a declining tendency contrary to the EC0Fy system. Meanwhile, the formation and enlargement of metalenriched interphase give rise to intensify ΔεInterphase of the EC5Fx system. According to previous studies, the ratio of

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CONCLUSIONS In summary, we exploited an accurate interfacial regulation strategy to realize percolation of mechanical performance toward striking elastomer reinforcement. Architectural design of interfacial metal−ligand bridges was implemented to establish firm anchoring between oxygenic moieties in the skeleton of ENR and surface functional moieties of CDs, leading to the formation of interfacial metal-enriched regions. The topological evolution of interfacial region was quantitatively determined by SAXS data. Specifically, the increase of metal−ligand bridges facilitated the progressive enlargement of the interfacial region which finally overlapped with neighboring domains, giving rise to a penetrating interphase. Subsequently, the volumetric enlargement of the interfacial region Φi exhibited a similar percolation behavior in sync with that of macroscopic mechanical performance. The relative dielectric strength, as another quantitative indicator of Φi, also confirmed the above conclusion. Furthermore, dielectric measurements revealed that the metal-enriched interphase performed much slower relaxation dynamics separated from the bulk polymer, which significantly improved the network rigidity in bulk and accounted for the prominent reinforcement of elastomer matrix. Hereto, the current system corroborates that the formation of penetrating interphase can be an executable mechanism to induce the reinforcement percolation of ENCs. We also envision that the implementation of interphase percolation mechanism can provide significant insights into the optimized design of high-performance elastomer materials.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.7b11239. Stress−strain and loss tangent curves of EC0Fy samples, X-ray scattering profiles of ENR samples, dielectric loss profiles for EC5F0 and EC0F1.5 samples, description of the KT-modified hard-sphere model, fitting parameters for KT-modified model, mechanical properties, VFT fitting parameters, and dielectric strength for ENR samples (PDF) H


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AUTHOR INFORMATION

Corresponding Author

*(B.G.) E-mail [email protected]; Tel +86 20 87113374; Fax +86 20 22236688. ORCID

Baochun Guo: 0000-0002-4734-1895 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by National Basic Research Program of China (2015CB654703), National Natural Science Foundation of China (51673065, 51703064, U1462116, and 51473050), and Natural Science Foundation of Guangdong Province (2014A030310435 and 2014A030311051).

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