Polymer 55 (2014) 5389e5395
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Mechanical properties of polybutadiene reinforced with octadecylamine modified graphene oxide Yan Zhang a, James E. Mark a, *, Yanwu Zhu b, Rodney S. Ruoff c, Dale W. Schaefer d, ** a
Department of Chemistry, University of Cincinnati, Cincinnati, OH 45221-0172, USA Department of Materials Science and Engineering and CAS Key Laboratory of Materials for Energy Conversion, University of Science and Technology of China, Hefei, Anhui 230026, China c Department of Mechanical Engineering and The Texas Materials Institute, University of Texas at Austin, Austin, TX 78712-0292, USA d Department of Biomedical, Chemical and Environmental Engineering, College of Engineering and Applied Science, University of Cincinnati, Cincinnati, OH 45221-0012, USA b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 June 2014 Received in revised form 18 August 2014 Accepted 24 August 2014 Available online 3 September 2014
Octadecylamine-modified graphene-oxide (OMGO) polybutadiene nanocomposites with different OMGO loadings were prepared by solution mixing. The dispersion of OMGO in chloroform is greatly improved compared to GO. Toughness and elongation of PBDeOMGO nanocomposites increase by 332% and 191% respectively compared with pure PBD. However, Young's modulus of PBDeOMGO nanocomposite decreases by 10% at 2-wt% loading. Graphene sheet crumpling accounts for the increased toughness, the absence of modulus reinforcement and the absence of a Payne effect for PBDeOMGO. The oxidation susceptibility of PBD is greatly reduced after the addition of OMGO, which is particularly desirable in the tire industry. © 2014 Elsevier Ltd. All rights reserved.
Keywords: Polybutadiene Graphene oxide Mechanical property
1. Introduction Rubbers (elastomers) are an important class of commercial polymers. Classic elastomers, such as polybutadiene and polyisoprene are used as general purpose rubbers in high volume products such as tires, hoses, belting, and flexible automotive parts [1]. Rubber materials have been extensively studied because of easy processing, flexibility and excellent thermal properties [2e4]. However, rubber is commonly used in form of composites since pure rubber lacks the required mechanical properties such as wear resistance and strength. Rubber nanocomposites are the class of filled rubbers in which at least one dimension of the fillers is on the nanometer scale. Most commonly used fillers are silica, carbon black, clay and carbon nanotubes [5e7]. Increased modulus is achieved at relative high filler loading such as 20e50 per hundred rubbers (phr), which can reduce toughness due to defects caused by the fillers [8]. However, the incorporation of small-size fillers in cross-linked elastomers results in specific nonlinear mechanical behaviors including Payne
* Corresponding author. Tel.: þ1 513 556 9292; fax: þ1 513 556 9239. ** Corresponding author. Tel.: þ1 513 556 5431; fax: þ1 206 600 3191. E-mail addresses:
[email protected] (J.E. Mark),
[email protected],
[email protected] (D.W. Schaefer). http://dx.doi.org/10.1016/j.polymer.2014.08.065 0032-3861/© 2014 Elsevier Ltd. All rights reserved.
effect and Mullins effect. The Payne effect is typically observed at small strain. The dynamic storage modulus decreases strongly with increasing strain amplitude [9]. Mullins et al. first reported that the degree of softening increases with increasing stiffening ability of the fillers [10]. Some Mullins softening is observed in carbon black and silica filled rubber composites systems [11]. Polybutadiene (PBD), a synthetic rubber, has a higher resistance to wear over styrene-butadiene rubber and natural rubber, which are its main competitors in rubber-industry applications due to their lower glass transition temperatures [12]. PBD is a low cost rubber used for soles, gasket, seals and belts [13]. PBD is normally formulated with fillers, such as silica or carbon black. Graphene is an emerging filler candidate that has been widely studied in thermoplastics, but not in elastomers. Graphene shows high thermal conductivity (5000 W m1 K1) [14e16], highest Young's modulus ever measured (1 TPa) [17] and large theoretical surface area (2675 m2 g1) [18]. High modulus and large surface area promise dramatic improvement in mechanical properties, which as yet has not been realized. Within the few published paper on elastomers filed with graphene materials, Araby et al. [19] reported that tensile strength of styrene butadiene rubber filled with graphene increases by 230% using melt compounding. However, in order to obtain such improvement, a large amount of graphene (24%) was incorporated into rubber, which causes defects in products and increases cost.
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An effective approach to achieve graphene-based polymer composites is based on chemical transformation of graphite to graphite-oxide (GO), which readily disperses in water and exfoliates to form individual, single-layer graphene oxide sheets [20e23]. Further modification of GO is necessary to achieve fully dispersed GO in common organic solvents [24e28]. Here we demonstrate a modification scheme based on octadecylamine (ODA) modification of GO that greatly improves GO dispersion in chloroform. We report the synthesis of ODA-modified GO (OMGO) and examine the mechanical properties of PBDeOMGO nanocomposites as a function of OMGO loading. The amino group in the ODA modifier reacts with carboxylic acid groups in GO. In contrast to the preponderance of graphene-based thermoplastic composites, we investigate the performance of thermosetting rubber materials. Toughness and elongation at break of PBDeOMGO improve by 332% and 343% respectively at 2-wt% OMGO. However, Young's modulus of PBDeOMGO decreases by 10% at OMGO loading 2 wt%. 2. Experimental section 2.1. Materials GO, prepared by a modified Hummers method [20]. ODA and N, N0 -diisopropylcarbodiimide (DIC) were purchased from SigmaeAldrich Co. Chloroform was purchased from Tedia Company Inc. Anhydrous acetonitrile was purchased from Acros Organics. High cis-1, 4-polybutadiene (c-PBD) was purchased from SigmaeAldrich Co. Dibenzoyl peroxide (BPO) was purchased from Acros Organics. All reagents were used as received. 2.2. Preparation of OMGO ODA-modified GO (OMGO) was made using the reaction between carboxylic acid groups and epoxy groups from GO and amino groups (using DIC to activate carboxylic groups), as shown in Scheme 1. The selectivity of carboxylic acid groups and epoxy groups remains an open question [29,30]. Briefly, the desired amount of GO was dispersed in 40 ml anhydrous acetonitrile followed by ultrasonication for 1 h. DIC was added into graphene oxide dispersion followed by reaction for 4 h at 70 C to activate the
Scheme 1. Synthesis of OMGO with octadecylamine.
carboxylic acid groups. Then ODA was added into the mixture followed by refluxing at 75 C overnight under nitrogen atmosphere. After the reaction, the OMGO was purified by washing with dimethyl formamide and acetone successively to remove residual DIC and unreacted amines. Black, solid OMGO was obtained after vacuum drying at 50 C overnight. 2.3. Preparation of the unfilled cis-PBD networks The desired amounts of polymer were first dissolved in chloroform. After a clear solution was obtained, 2-wt% BPO was added [31]. The solution was stirred at room temperature for 2 h and then transferred into Teflon™ dishes that were covered with aluminum foil for overnight solvent evaporation. Films were pressed at approximately 1.2 104 psi, 130 C for 2 h. The final samples were approximately 1.0 mm thick. 2.4. Preparation of PBDeOMGO nanocomposites The PBDeOMGO nanocomposites were prepared with various loadings of OMGO. Firstly OMGO was dispersed in chloroform with the aid of ultra-sonication for 1 h to yield a well-dispersed solution. Secondly, the well-dispersed OMGO solution was mixed with c-PBD with 0.50, 1.00 and 2.00 weight ratio for 2 h following the same procedures as for unfilled c-PBD discussed above. 2.5. Characterization Fourier transform infrared spectroscopy (FTIR) recorded on a Nicolet 6700 (Thermo Scientific) spectrometer was used to characterize the chemical structure of GO and OMGO. Samples were measured under a mechanical force by pressing the un-exfoliated sample's surface against a diamond window. FTIR spectra were collected in the range 4000e450 cm1. Thermo-gravimetric analyses (TGA) were done on the GO and OMGO powders using a TA Q500 instrument (TA Instruments) under a nitrogen atmosphere with a heating rate of 5 C min1. Degradation temperatures of PBDeOMGO nanocomposites were measured using NETZSCH STA 409 instrument under a purge flow of 20 ml/min argon at heating rate of 20 C min1 from 25 to 700 C. Powder X-ray diffraction (XRD) measurements were carried out using a PANalytical X'Pert Pro MPD diffractometer with Cu Ka radiation (l ¼ 1.541 Å) at 45 kV and 40 mA. The diffraction angle was increased from 5 to 35 with the scanning rate of 0.05 min1. Scanning electron microscopy (SEM) was conducted with an FEI Phillips Electroscan XL30 ESEMeFEG microscope using an acceleration voltage of 15 kV. OMGO suspensions (0.03 mg ml1) were spin-coated onto a flat aluminum plate at 2000 r.p.m. for 30 s. Then the OMGO-coated aluminum plate was mounted on a standard specimen holder using a double-sided carbon conductive tape for SEM imaging. A transmission electron microscopy (TEM) sample was prepared by placing a few drops of dispersion onto a lacey carbon film support on a Cu grid. Images were acquired in a JEOL 1230 transmission electron microscope operated at 80 kV. Atomic force microscopy (AFM) images were obtained using a Dimension 3100 AFM made by Veeco Instruments Inc., operated in tapping mode using Veeco RTESP type silicon cantilevers with a resonance of frequency of 360 kHz. The samples for AFM measurements were prepared by ultrasonic treatment of OMGO in chloroform for 1 h, followed by spin-coating OMGO suspensions in chloroform (0.03 mg ml1) on freshly cleaved mica surfaces at 2000 r.p.m. for 30 s and then drying under vacuum at room temperature. The tensile properties were measured using an ESM-301 (MARK-10) tensile tester. The experiments were carried out at
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room temperature at a crosshead speed of 50 mm/min using dumbbell-shaped specimens with an original length of 40 mm. Three dumbbell-shaped samples were measured and standard deviation of three measurements is the error bar. MettlereToledo DMA-861 was used to measure the viscoelastic data in order to investigate the Payne effect of nanocomposites. The tests were performed on rectangular samples in shear mode at 5 Hz at different shear amplitude. 3. Results and discussions FTIR data shows that ODA is linked to the GO by the reaction with carboxylic acid group. Fig. 1 shows the FTIR spectra of GO and OMGO. In the spectrum of GO, broad peak between 3550 and 2500 cm1 (OeH stretching from COOH and OH groups), C]O (1714 cm1), and aromatic C]C (1595 cm1) stretches were observed. After modification with ODA, there is a peak between 3500 and 3000 cm1, which is attributed to OeH stretching from OH and two new peaks appear at around 2914 cm1 and 2839 cm1, which are assigned to the CeH stretch in the methylene group from ODA [32]. The peak at 1465 cm1 is due to the asymmetrical bending vibration of methyl group from ODA. The reaction of carboxylic acid groups with amine groups in the ODA is confirmed by the observed decrease in wavenumber for the peak of C]O stretch, from 1714 to 1570 cm1, which overlaps with C]C peak. Since the intensity of peak is proportional to number of functional groups in the system, the peak at 1570 cm1 is relatively weak. TGA was utilized to monitor the decomposition of functional groups. Fig. 2 shows the typical TGA curves of GO and OMGO. In the GO curve, relatively low thermal stability was observed. GO starts losing mass below 100 C (about 7%), which is attributed to the absorbed water. The 46% weight loss between 100 C and 280 C is due to decomposition of the labile oxygen-containing functional groups. Between 280 C and 600 C, no obvious weight loss was observed indicating all of the oxygen functional groups have decomposed below 280 C. The decrease in mass around 600 C is due to the pyrolysis of the carbon skeleton of the GO. After modification with ODA, no weight loss is observed in OMGO below 100 C and the weight loss between 100 and 280 C is 9 wt% which is 37 wt% lower than that of GO (46 wt%). The 37 wt% difference is
Fig. 1. FTeIR of GO and OMGO. The peaks at 2914 cm1 and 2829 cm1 appear after modification with ODA assigned to CeH stretch in the methylene group from ODA, which indicates ODA was attached to GO.
Fig. 2. TGA curves of GO and OMGO. GO shows 37% more weight loss compared to OMGO between 100 and 280 C, indicating COOH groups have been converted after ODA modification.
attributed to ODA modifier, which degrades above 280 C. We conclude that 37 wt% of oxygen functional groups in GO have been modified by ODA [33,34]. This value is consistent with the weight loss (38 ± 2 wt%) between 280 C and 550 C, which is assigned to the removal of alkyl chains. The consistency of these weight losses confirms the mass loss between 100 and 280 C is due to ODA modifier. XRD was used to determine the morphology and the layer spacing of dry GO and OMGO. Fig. 3 shows a diffraction peak at 2q ¼ 10.3 indicating an interlayer distance of 0.86 nm for GO. After modifying GO with ODA, the 10.3 peak almost disappears. However, a broad peak appears with comparable intensity at 22 , which is closer to the typical (002) diffraction peak of graphite (26.6 ), corresponding to 0.40 nm of interlayer distance. The same 22 peak is reported in the work of Dubin et al. [35] and Pei et al. [36] Compared with diffraction peak of graphite, which is extremely
Fig. 3. XRD pattern of GO and OMGO. After modifying GO with ODA, the peak at 10.3 disappears. The broad peak of OMGO between 15 and 30 indicates that OMGO sheets are disordered stacks.
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Fig. 4. (A): SEM image of OMGO indicating OMGO shows a layered structure; (B): TEM image of OMGO. Extremely thin layers were observed indicating that OMGO was fully dispersed in chloroform.
Fig. 5. AFM images and height profiles of OMGO dispersed at a concentration of 0.03 mg ml1 in chloroform. 5.0 mm 5.0 mm scan area was selected to measure the thickness of OMGO, which is 0.9 nm.
sharp and intense, the broad peak of OMGO between 15 and 30 indicates that OMGO sheets are disordered stacking. Under sonication OMGO sheets easily separate into single layers as shown in Fig. 4B. SEM and TEM measurements were performed to investigate the morphology of OMGO. Fig. 4A shows the SEM of OMGO at 500 nm scale indicating that OMGO displays a layered structure. The OMGO observed with TEM (Fig. 4B) is in the form of flat sheet. In the darker part on the right corner is carbon grid. Most of the area on the left is light gray color, which is covered by very thin layer of dispersed OMGO. Fully exfoliated OMGO sheets were achieved after ODA modification. Some areas are darker, which indicates stacking of the OMGO sheets. AFM was used to determine the thickness of dispersed OMGO (Fig. 5). The cross-sectional view of the AFM image of OMGO indicates that the thickness of OMGO is 0.9 nm. The typical observed monolayer GO sheet is 0.8 nm, which is larger than the theoretical graphite sheet with van der Waals thickness of ca. 0.34 nm because of the presence of oxygen functional group above and below the GO plane. So our result indicates that the OMGO platelet is a single layer, which can also be seen from TEM image. TGA was employed to examine the degradation of PBDeOMGO nanocomposites. Fig. 6 is a composite plot of the TGA mass loss curves obtained from the 0, 0.5, 1 and 2 wt% OMGO filled systems. The introduction of OMGO into PBD increases the degradation
Fig. 6. TGA of PBDeOMGO nanocomposite with different filler loading. At 2 wt% OMGO, the onset temperature of the PBDeOMGO composites is 5 C higher than pure PBD.
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Table 1 Degradation temperature of PBDeOMGO nanocomposites from Fig. 6. PBDeOMGO nanocomposites Loading Td ( C)
Pure PBD 380
0.5 phr 381
1.0 phr 385
2.0 phr 385
temperature (Fig. 6). The derivative of each mass loss curve defines the degradation onset temperature as shown in Table 1. At 2 wt% OMGO, the onset temperature of the PBDeOMGO composites is 5 C higher than pure PBD. Different aspects of the mechanical properties of PBDeOMGO nanocomposites were examined. Toughness greatly improves while the Young's modulus of PBDeOMGO nanocomposites decreases slightly. Fig. 7A shows the stressestrain curves of PBDeOMGO nanocomposites as a function of filler loading. The areas under the curves are the toughnesses (Fig. 7B) and the maximum extension reflects the elongation of composites (Fig. 7C). The toughness and elongation increase 332% and 191% respectively at filler loading of 2.0 wt%. The slope of the curves in the initial small strain portion (Fig. 7D) is Young's modulus. The Young's modulus decreases, which indicates the stiffness of the materials does not increase by adding OMGO as filler. The fact that Young's modulus decreases and toughness increases by the introduction of OMGO is due to single-sheet nature of OMGO, which adopts a wrinkled morphology that imparts a measure of entropic, rubberlike elasticity to the filler itself [37,38]. That is, because the bending modulus of the sheets is so low, adding filler does not enhance the modulus. In fact, the filler may interfere with crosslinking chemistry leading to decreased modulus. On the other hand, stretching of the crumpled OMGO sheet adds an extra
Fig. 8. Storage modulus (E0 ) aWs a function of shear amplitude at 5 Hz. The data show almost no Payne effect (reduction of modulus with strain amplitude). These data indicate that the crumpled OMGO has a lower modulus than the matrix.
reinforcement mechanism at large elongations, which improves toughness and elongation at break. The samples were tested using dynamic mechanical analysis (DMA) to elucidate the cause of the low Young's modulus. A typical DMA result on PBDeOMGO nanocomposites deformed in shear mode is shown in Fig. 8. Pure PBD has a larger elastic modulus than
Fig. 7. A. Stressestrain curves of PBDeOMGO composites with different filler loading. Fig. 7B. Toughness of PBDeOMGO naocomposites vs filler contents Fig. 7C. Elongation of PBDeOMGO nanocomposites vs. filler loading. Fig. 7D. Young's modulus of PBDeOMGO composites vs. filler loading.
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Fig. 9. Comparison of the Payne effect for PBDeOMGO and CNFePU. The storage modulus for PBDeOMGO is calculated from shear modulus in Fig. 8 assuming a Poisson ration of 0.5. CNF reinforces PU (and shows a Payne effect) even though PU matrix has a higher zerostrain modulus than PBD. PBD shows neither reinforcement nor a Payne effect.
PBDeOMGO. The modulus decreases as filler loading increases and reaches minimum at 1.0 wt%, which matches the tensile result (Fig. 7D). Within error, there is no observable Payne effect (decrease of modulus with strain amplitude). The Payne effect is necessarily absent in PBDeOMGO because OMGO does not reinforce PBD. Fig. 9, compares the straindependence of the storage modulus of PBDeOMGO with that of carbon-nanofiber (CNF) reinforced polyurethane (PU) [38], which does show a Payne effect. The latter reinforces PU at low strain, but at strain amplitude of ~10%, the modulus of the matrix is recovered. For OMGO the modulus of the matrix is already “recovered” at zero strain so it cannot display a Payne effect. The Payne effect is generally attributed to clustering of the filler particles induced by van der Waals forces [39]. Clustering is disrupted by strain, leading to reduced modulus enhancement and, at high strain, no modulus enhancement (the Payne effect). Apparently the crumpled morphology of OMGO precludes the formation of deck-of-cards configurations required for strong van der Waals attractions. In the case of CNFs, on the other hand, side-by-side fiber configurations are present, at least over the persistence length of the fibers [37,38], which leads to the clustering required for both zero-strain modulus enhancement and the Payne effect. After the tensile test, the fracture surfaces were examined by SEM. Fig. 10A and B shows the cross-section SEM images at fracture surface of pure PBD and PBDeOMGO with 0.5 wt% OMGO. No
Fig. 10. (A) Cross-section SEM image of pure PBD at the fracture surface; (B) Cross-section SEM image of PBDeOMGO with 0.5 wt% OMGO at the fracture surface. No aggregation is observed in the PBD matrix, which indicates that OMGO is fully exfoliated.
Fig. 11. (A): Pure PBD; (B): 0.5 wt% OMGO; (C): 1 wt% OMGO; (D): 2 wt% OMGO after 12 months.
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aggregation is observed in the PBD matrix after the addition of OMGO, which indicates that OMGO is fully exfoliated. The degradation of PBDeOMGO nanocomposites was observed at times up to 12 months after sample prepared. Fig. 11 compares pure PBD with PBD filled with OMGO at three different percentages after 12 months. PBD has a double bond in the backbone, which is easily oxidized by ozone when exposed to air [40]. After 12 month, pure PBD becomes yellow, brittle and forms cracks. The addition of OMGO greatly reduces oxidation. As the OMGO loading increases, PBDeOMGO nanocomposites remain flexible. No cracking is observed for 1 wt% and 2 wt% OMGO. The absence of oxidation in PBDeOMGO nanocomposites might be due to the filler-polymer attraction. OMGO has many C]C moities, which can form pep interactions with C]C in PBD structure. This observation is very important to industry since PBD is the most important rubber component in the tire industry. 4. Conclusions GO was successfully modified by ODA. ODA attaches to surface by reaction of surface carboxylic acid groups and epoxy groups on GO with the amine group on ODA. The introduction of ODA onto GO greatly improves exfoliation in chloroform and facilitates casting of PBD composites. OMGO in the PBDeOMGO nanocomposites acts as a spring, which stores energy leading to the 332% improvement in toughness and 191% improvement in elongation. However, Young's modulus of PBDeOMGO nanocomposites decreases, indicating the crumpled sheets have a lower modulus than the rubber itself. Since there is no modulus enhancement there is also no filler-induced the Payne effect. The oxidation of PBD decreases after the addition of OMGO. Acknowledgments This work was supported by the National Science Foundation through Grant DMR-0803454 (Polymers Program, Division of Materials Research). We thank the Advanced Materials Characterizing Center (AMCC) at University of Cincinnati for providing the instruments and Professor Vesselin Shanov's help on tensile measurements. References [1] El Fray M, Goettler LA. Application of rubber nanocomposites. In: Rubber nanocomposites. John Wiley & Sons, Ltd; 2010. p. 675e96. [2] Wang J, Vincent J, Quarles CA. Review of positron annihilation spectroscopy studies of rubber with carbon black filler. Nucl Instrum Methods Phys Res Sect B 2005;241(1e4):271e5. [3] Chen G, Zhao W. Rubber/graphite nanocomposites. In: Rubber nanocomposites. John Wiley & Sons, Ltd; 2010. p. 527e50. [4] Ma J, Zhang L-Q, Geng L. Manufacturing techniques of rubber nanocomposites. In: Rubber nanocomposites. John Wiley & Sons, Ltd; 2010. p. 21e62. [5] Rohini Thimmaiah S, Siddaramaiah. Investigation of carbon black and metakaolin cofillers content on mechanical and thermal behaviors of natural rubber compounds. J Elastomers Plast 2013;45(2):187e98. [6] Song M, Wong C, Jin J, Ansarifar A, Zhang Z, Richardson M. Preparation and characterization of poly(styrene-co-butadiene) and polybutadiene rubber/clay nanocomposites. Polym Int 2005;54(3):560e8. [7] Zhou D, Mark JE. Preparation and characterization of trans-1,4-polybutadiene nanocomposites containing in situ generated silica. J Macromol Sci Pure Appl Chem 2004;41(11):1221e32. [8] Patwardhan S, Taori V, Hassan M, Agashe N, Franklin J, Beaucage G. An investigation of the properties of poly(dimethylsiloxane)-bioinspired silica hybrids. Eur Polym J 2006;42(1):167e78. [9] Chazeau L, Gauthier C, Chenal JM. Mechanical properties of rubber nanocomposites: how, why … and then?. In: Rubber nanocomposites. John Wiley & Sons, Ltd; 2010. p. 291e330.
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