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bismaleimide composites by using reduced graphene oxide with non-covalent functionalization
Accepted Manuscript Enhanced mechanical and tribological properties of graphene/bismaleimide composites by using reduced graphene oxide with non-covalent functionalization Chao Liu, Yufei Dong, Yang Lin, Hongxia Yan, Wenbo Zhang, Yan Bao, Jianzhong Ma PII:
S1359-8368(18)33035-X
DOI:
https://doi.org/10.1016/j.compositesb.2019.02.005
Reference:
JCOMB 6591
To appear in:
Composites Part B
Received Date: 14 September 2018 Revised Date:
25 December 2018
Accepted Date: 5 February 2019
Please cite this article as: Liu C, Dong Y, Lin Y, Yan H, Zhang W, Bao Y, Ma J, Enhanced mechanical and tribological properties of graphene/bismaleimide composites by using reduced graphene oxide with non-covalent functionalization, Composites Part B (2019), doi: https://doi.org/10.1016/ j.compositesb.2019.02.005. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Graphical Abstract (Figure)
ACCEPTED MANUSCRIPT Enhanced mechanical and tribological properties of graphene/bismaleimide composites by using reduced graphene oxide with non-covalent functionalization† Chao Liu a∗, Yufei Dong b, Yang Lin a, Hongxia Yan c, Wenbo Zhang a, Yan Bao b, Jianzhong Ma b
University of Science &Technology, Xi’an 710021, China
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a Shaanxi Collaborative Innovation Center of Industrial Auxiliary Chemistry and Technology, Shaanxi
b College of Bioresources Chemical and Materials Engineering, Shaanxi University of Science & Technology, Xi’an 710021, PR China
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c Key Laboratory of Space Applied Physics and Chemistry, Ministry of Education, Department of Applied Chemistry, School of Science, Northwestern Polytechnical University, Xi’an 710129, China ABSTRACT: In order to simultaneously improve the dispersibility of graphene in the bismaleimide
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(BMI) resin and the interfacial interaction between them, the polytriazine@RGO nanocomposite (PTZ@RGO) with terminal amine was obtained by the method of one-step precipitation polymerization. The PTZ@RGO hybrid nanocomposite was then chosen as filler to improve the mechanical and tribological properties of BMI resin. Through effective surface modification, the compatibility and interface strength between graphene and BMI matrix were significantly improved, leading to the flexural
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strength, impact strength, anti-friction, wear resistance and thermal stability of PTZ@RGO/BMI composites were significantly improved as compared to those of neat BMI resin and RGO/BMI composites. The other properties including hardness and thermal properties of composites were all
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improved with increasing amount of PTZ@RGO fillers. These high performance BMI composites have a great potential for application in aerospace, mechanical engineering and other automobile brake devices.
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Key words: Polymer-matrix composites (PMCs); Theromosetting resin; Mechanical properties; Wear.
1. Introduction
Bismaleimide (BMI) is a high-performance thermosetting polymer with outstanding thermal stability, high specific modulus, and upstanding mechanical properties under the condition of high temperature [1-4]. Meanwhile, it has the similar formability to that of epoxy resin. While, the friction coefficient and wear rate of it are higher when it is used as friction material. Which are the major causes for material and ∗
Corresponding author. E-mail address: [email protected] (Chao Liu) 1
ACCEPTED MANUSCRIPT energy losses in mechanical processes. In order to improve mechanical durability and energy efficiency, BMI needs to be modified to enhance its performance, especially the mechanical and tribological comprehensive properties [5-7]. Recently, in order to overcome those shortcomings, varieties lubrication fillers such as carbon
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nanotube, MoS2, graphene or their derivative were introduced into the BMI matrix by researchers [8-10]. Graphene, the 2-dimensional carbon allotrope with outstanding self-lubricating and physicochemical properties, has attracted great interest as self-lubricating solid or as additive for polymer based friction materials [11-13]. The introduction of graphene into polymer matrix can not only improve its tribological
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properties but also improve its mechanical properties [14-16]. However, neat graphene contains no surface functional groups and has very limited dispersibility in polymer matrix, seriously limiting its
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potential application in preparation of functional nanocomposites. The surface modification of graphene by grafting of polymers, and the applications of such polymer-grafted nanoparticles have become an area of major interest lately. We have reported various hyperbranched polymer chemically surface modified graphene, which is of great value for improving compatibility with polymer matrix, resulting in the excellent distribution and enhanced interfacial bonding between graphene and polymer matrix [17, 18].
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However, this method has many steps and the yield of the product is low. Meanwhile the modification of covalent bond can result in the destruction of graphene structure, thus reducing its performance. Recently, the method of fabricating the functionalized graphene by deposition of nanoparticles on graphene to
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prevent the overlay of graphene sheets and enhance its tribological properties has attracted increasing attention [19-21]. Nevertheless, an appropriate interface interaction is essential for effective interface load transfer and wear resistance reinforcements. Furthermore, the mobility of both graphene is likewise
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crucial for achieving toughening and lubricating properties of composites. To resolve these contradictions, one expects to design a rational interface structure to optimize the mechanical and tribological properties of graphene-polymer composites [22, 23]. In this contribution, the reduced graphene oxide (RGO), cyanuric chloride (CC) and hexamethylenediamine (HMD) are used to prepare the polytriazine@RGO nanocomposite (PTZ@RGO) with terminal amine by the method of one-step precipitation polymerization. The PTZ@RGO was incorporated into the BMI matrix to form PTZ@RGO/BMI binary composite (See Fig. 1). This method can not only effectively resolve the aggregation of graphene but also improve the compatibility and 2
ACCEPTED MANUSCRIPT interfacial adhesion between graphene and the polymer matrix. The effects of RGO and PTZ@RGO content on the mechanical and tribological properties of BMI matrix were investigated. The wear mechanical of the composites was discussed. 2. Experimental
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2.1 Reagents and materials Natural graphite flakes (325 mesh) was obtained from Qingdao Hensen Graphite Co., Ltd. The graphene oxide (GO) nanosheets were produced from natural graphite flakes by the modified Hummers method [24]. Cyanuric chloride (CC, 99%), hexamethylene-diamine (HMD), NaOH and triethylamine
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(TEA) were purchased from Aladdin Chemistry Co., Ltd. N, N-Dimethylformamide (DMF), hydrazine hydrate and anhydrous alcohol were purchased from Tianjin Tianda Chemical Co. Ltd. The 4,
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4`-Bismaleimidodiphenylmathane (BDM) was provided by Rongchang Ning research group at Northwestern Polytechnical University. Diallyl Bisphenol A (DBA) was purchased from Sigma-Aldrich. All reagents were of analytical grade and used as received without further purification. 2.2 Procedures
Typically, graphene oxide (GO) was synthesized from natural graphite flakes by a modified Hummers
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method [24]. For synthesizing PTZ@RGO nanocomposite: firstly, 0.2 g GO was dispersed in 200 mL deionized water and sonicated in an ultrasonic bath for 0.5 h. Then, 2 mL hydrazine hydrate was added into the above solution and the mixture was stirred constantly at 98 oC kept for 6 h to prepare the RGO.
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Secondly, the RGO (0.2 g), CC (0.8 g, 0.5 mol) and a small amount of NaOH were added into a 250 mL three-neck round bottom flask. A mixture of DMF (150 mL) and 2.0 mL of TEA as an acid-acceptor were subsequently added. After ultrasonic treatment (150 W) for 30 min, HMD (1.6 g, 1.7 mol) was then added.
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The solution was maintained at 85 °C for 3 h with the protection of nitrogen gas by ultrasonic treatment (150 W). The product was filtered and washed by ethanol several times, and dried at 60 °C in a vacuum overnight.
The BMI-based composites were prepared by mixing the pre-weighed quantities of nanofillers,
Diallyl Bisphenol A (DBA) and BDM (DBA and BDM with a mass ratio of 3:4). Then the mixture was heated to 135 °C till the mixture totally melting and the nanofillers dispersed uniformly. Finally, the mixture was cured following the schedule of 150 °C/2 h + 180 °C/2 h + 220 °C/4 h. A post curing process was 250 °C/4 h. The BMI resin-based materials were also prepared in similar procedures to perform 3
ACCEPTED MANUSCRIPT contrast experiments. 2.3 Characterization Flexural strength and impact strength were performed according to GB/T2567-2008. The samples dimension for impact test and flexural test were (80 ± 0.2) × (10 ± 0.2) × (4.0 ± 0.2) and (80 ± 0.2) × (15
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± 0.2) × (4.0 ± 0.2) mm3, respectively. Each value obtained represented the average of five samples. The friction and wear tests were conducted on a test machine (MM-200, load 196 N, speed 200 rpm/min) under dry condition (temperature: 20±5oC, humidity: 50±10%) according to the GBT/3960-1983 (Chinese Standard). Where the pin and friction pair are the composites specimen and
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#45 steel disk, respectively. Before each test, the counterpart steel ring and the materials were sanded using No. 600 water-abrasive paper. Then the steel ring and samples were cleaned by acetone. The
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calculation of volume wear rate can be detailed in supporting information 1.
The Fourier transform infrared (FT-IR) spectra of samples was carried obtained with a Bruker IFS66/S FT-IR spectrometer. XRD patterns of the powered were measured on a Philips X’Pert MPD system with Cu-Ka radiation. X-ray photoelectron spectra (XPS) were investigated using a PHI Quantum 2000 Scanning ESCA Microprobe system. All XPS spectra were corrected using the C1s line at 284.6 eV.
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Raman spectra were obtained carried out on a Horiba Scientific XploRa Raman spectrometer. The microstructures of the composite specimens were characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM) via a Hitachi S-4800 and a FEI Tecnai G2 F20 microscopy,
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respectively. The hardness of the composite specimens was performed on a MH-5 Vector microhardness meter at a load of 50 N and a loading time of 20 s. Thermogravimetric analysis (TGA) was conducted in TGA Q50 at a heating rate of 10 oC/min, in a nitrogen atmosphere. Friction final temperature of the
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composites friction surface was recorded by a TMP-2 multichannel temperature tester. 3. Results and Discussion
3.1 Characterization of PTZ@RGO nanocomposite The FT-IR spectra of GO and PTZ@RGO are shown in Fig. 2. For the GO, the absorption peaks of GO appears at 1056, 1638 and 3430 cm-1, which are attributed to the stretching vibration of C=C, C-O and -OH groups (or the physical adsorption of water) on GO surface, respectively [25]. In the spectrum of PTZ@RGO, the peaks at 1443 and 1551 cm-1 are assigned to the skeleton vibration of the triazine ring, respectively. Additionally, the peak at 1330cm-1 is due to the C-N stretching vibration of the carbon atoms 4
ACCEPTED MANUSCRIPT of triazine rings linked with the -NH-. However, there is only one absorption peak at 3497 cm-1 between 3400 and 3500 cm-1, which may be attributed to the overlapping of the absorption peak of water and -NH2. The series of the peaks suggest the existence of PTZ in the PTZ@RGO nanocomposite [26]. XRD analysis was performed to investigate the structures of the samples. Fig. 3 shows the XRD
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patterns of GO, RGO and PTZ@RGO. As shown in Fig. 3, the XRD spectrum of the GO shows one strong (001) characteristic peak at 11.3° with d-spacing of 0.78 nm. After being reduced, the d-spacing of GO is reduced from 0.78 to 0.36 nm, which confirms that most oxygenic groups are removed. The XRD spectrum of RGO shows a very broad diffraction peak at 24.3° and the diffraction peak at 11.3°
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disappears, confirming the formation of RGO [27]. Compared with the XRD pattern of RGO, no obvious diffraction peaks for PTZ@RGO have been observed, which may be due to the formation of the
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PTZ@RGO nanocomposite so that individual RGO remains isolated [28].
XPS characterization was employed to analyze the surface chemical compositions of RGO and PTZ@RGO. In Fig. 4a, the peaks at 285.1, 399.6 and 532.0 eV are attributed to Cl2p, C1s, N1s and O1s, respectively [29]. From the C1s spectrum shown in Fig. 4b, four peaks attributed to C=C at about 284.5 eV, C-OH at about 286.6 eV, C-O-C at about 287.5 eV, and C=O at about 288.7 eV were observed. As
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presented in Fig. 4c, the C1s XPS peak-fitting of PTZ@RGO shows six different peaks, centered at 284.5, 285.6, 287.5, 288.1 and 288.7 eV, corresponding to C-C, C-N, C-O-C, C=N and C=O groups, respectively. However, the peaks at 287.5 and 288.7 eV have almost disappeared, which confirms a considerable
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degree of reduction. The peaks at 285.6 and 288.1 eV confirm that the existence of PTZ [30]. These results are in good agreement with the results of FT-IR. In order to further investigate the morphologies of the prepared hybrids, HR-TEM images were
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recorded. As shown in Fig. 5a and 5b, the GO and RGO have a typical shape resembling the wrinkled thin paper flake, and each nanosheet is up to several microns in size. After depositing PTZ onto the surface of RGO, a significant change can be observed (See Fig. 5c). The image of PTZ@RGO nanosheets show a gray texture of organic materials uniformly covering the entire surface of graphene nanosheets. This is attributed to that the polytriazine framework is rich in conjugated structure, which makes it can be firmly anchored to the surface of graphene [31]. These TEM observations are consistent with the other results from the other characterization methods, again demonstrating the successful deposition of PTZ to the graphene surface. The PTZ@RGO also shows excellent thermal stability (See Fig. S1). 5
ACCEPTED MANUSCRIPT 3.2 Characterization of PTZ@RGO/BMI composites When served in applications with high load and sliding velocity, it is of great importance for the composites to have outstanding flexural and impact strength to resist the deformation and bending failure. Based on this consideration, the flexural strength of BMI composites were also investigated. It can be
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clearly seen that the composites containing the PTZ@RGO exhibit higher strength values than those of the RGO at almost all the filler content (See Fig. 6a). The flexural strength of RGO/BMI increases continuously with the addition of RGO and reaches a maximum value of 146.8 MPa at 0.4 wt%. For 0.4
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wt% PTZ@RGO content, the flexural strength of its composites increased to 167.5 MPa, which is increased as much as 26.5% in comparison with that of neat BMI resin (132.4 MPa). This phenomenon is attributed to the unique sheets structure of hybrid graphene and the interfacial adhesion between
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PTZ@RGO and BMI matrix, which can greatly improve the flexural strength of BMI resin [31]. Fig. 6b shows the effect of the loading concentration of fillers on the impact strength of BMI composites. It shows that the suitable amount of fillers can properly improve the impact strength of BMI resin. The impact strength of the BMI composites respectively incorporating 0.6 wt% RGO and PTZ@RGO reaches the maximum value of 14.84 and 17.36 kJ/m2, increasing as much as 25.0 % and
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46.3 % compared to that of neat BMI resin (11.87 kJ/m2). Unfortunately, the flexural and impact strength of the composites decline when the amount of fillers is further increased, which may be due to the fact that excessive fillers can not be well dispersed in the BMI matrix and agglomerate to cluster [32].
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Consequently, the advantages of the fillers can not be fully utilized and the mechanical properties of the composites decrease with the uneven distribution of the fillers in the BMI matrix [33].
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To further research the fracture characteristics of composites, SEM images of the facture surfaces of neat BMI and its composites after impact testing are shown in Fig. 7. The crack propagation regions of neat BMI resin and its composites are different. The fracture surface of neat BMI resin showed in Fig. 7a is comparatively slippery, which exhibits a typical brittle feature. After the introduction of fillers, significant difference on the facture surfaces of composites is shown in Fig. 7b, 7c and 7d, respectively. The RGO is obviously protruded and aggregated on the fracture surface of 0.6 wt% RGO/BMI composite, which is indicative of poor compatibility and weak interfacial interaction between the RGO and BMI matrix (See Fig. 7b). As described in Fig. 7c, the fracture surface of 0.6 wt% PTZ@RGO/BMI composite is indented and containing many ductile sunken areas, indicating a typical tough fracture features, which 6
ACCEPTED MANUSCRIPT is considered from a result of the uniform distribution of PTZ@RGO in the BMI matrix and good interface interaction between PTZ@RGO and BMI matrix. The results are consistent with those obtained by TEM (See Fig. S2). However, when the content of PTZ@RGO is further increased, the fillers agglomerated in the matrix (See Fig. 7d), and induced a deterioration of the mechanical properties.
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There are two main differences between RGO and PTZ@RGO, which create the different mechanical properties of the composites. On the one hand, the existence of the abundant polar functional groups on the PTZ@RGO surface gives rise to a homogenous distribution of PTZ@RGO in BMI matrix. This state of dispersion probably results in enhancement of the mechanical interlocking between
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PTZ@RGO and BMI polymer chains [34, 35]. These strong interfacial interactions enable graphene to bring its amazing mechanical properties into resin matrix, achieving a perfect reinforcing and toughening
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material [36]. On the other hand, compared with the relatively inert surface chemistry of RGO, PTZ@RGO owns abundant amino groups which can form covalent bonds with the carbon-carbon double bonds of BMI (The reaction mechanism of PTZ and BMI is shown in Fig. S3). The corresponding mechanism of the reaction has been discussed in our previous studies [37]. These chemical bonds contribute to the efficient load transfer between matrix and fillers, and play an excellent role in inhibiting
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crack propagation, thus resulting in higher impact strength [38]. And the good interface interaction between PTZ@RGO and BMI matrix can enhance the wear-resisting ability of the composites. Fig. 8a shows the friction coefficients of composites as a function of fillers content for steady-state
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sliding against the counterpart steel ring under dry conditions. It can be seen that, the friction coefficient of composites decreases significantly with the addition of fillers. When the content of RGO is 0.6 wt%, the friction coefficients of RGO/BMI composites reach the minimum value (0.26).The friction coefficient
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of PTZ@RGO/BMI composites reach the minimum value of 0.20 at 0.6 wt% PTZ@RGO, decreasing as much as 42.9% in comparison with that of neat BMI resin (0.35). This may be due to the fact that the fillers in the resin matrix near the worn surface is exposed and acted as the lubricating thin “carbon film” on the worn surface of the steel counterpart during the wear process. As the content of fillers exceeds 0.6 wt%, the friction coefficient of composites increases. But the friction coefficient of the composites is still lower than that of the neat BMI resin. This phenomenon can be explained that the excessive fillers agglomerates in the resin matrix and thus leads to more micro-cracks on the composites surface, and therefore reduces the surface smoothness of composites during the wear process [39]. At the same time, 7
ACCEPTED MANUSCRIPT the friction coefficient of PTZ@RGO/BMI composites is more stable than those of RGO/BMI composites
in the process of friction (See Fig. S4). This may be attributed to that the modification can improve the dispersibility of graphene in BMI resin, thus uniform self-lubricating transfer film can be formed on the surface of counterpart steel ring during the process of friction.
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Fig. 8b shows the volume wear rate of composites with different fillers content. It can be seen that the volume wear rate of composites decreases initially and then increased with the increasing of fillers. When the content of RGO is 0.6 wt%, the volume wear rate of the RGO/BMI composites reaches the lowest value by 4.12×10-6 mm3/ (N•m). The volume wear rate of PTZ@RGO/BMI composites also
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reaches the lowest value (1.72×10-6 mm3/ (N•m)) at the same fillers content, decreases as much as 89.2 % compared to that of the neat BMI resin (15.94×10-6 mm3/ (N•m)). This result is in agreement with the test
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results of impact strength, which mainly attributed to composites containing 0.6wt％ fillers can absorb more energy and increase the load-carrying capability of the composites, thus the deformation of the composites could be reduced effectively during the wear process [40]. However, the volume wear rates of composites slightly increase when the fillers content exceeds 0.6 wt%. This may be due to that excessive
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fillers agglomerate in the BMI matrix, thus reduces the load-carrying capability of the composites during the wear process [41]. The tribological properties of 0.6 wt% PTZ@RGO/BMI composite under different loads (See Fig. S5) and sliding speed (See Fig. S6) are also investigate. In order to further investigate the wear mechanism of the composites under dry friction condition,
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the worn surfaces morphologies of different types of composites and corresponding surfaces of counterpart steel ring were examined by SEM, respectively. For neat BMI (See Fig. 9a), the worn surface
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of it is rough and some loose plate-like debris can be observed. On the surface of counterpart steel ring (See Fig. 9d), the transfer film is thick and discontinuous and piles of furrow can be seen. These phenomena reflect that the neat BMI has relatively poor wear resistance in its sliding against the steel counterpart and its tribological mechanism mainly follows adhesive wear and fatigue wear mechanism. As shown in Fig. 9b, the worn surface of 0.6 wt% RGO/BMI composite is compact and less apt to form debris from the matrix. The uneven areas on the worn surface of RGO/BMI composites and the discontinuous transfer film on the surface of counterpart steel ring (See Fig. 9e) may be attributed to the non-uniform distribution of RGO in the BMI matrix. Fig. 9c shows the worn surface of 0.6 wt% PTZ@RGO/BMI composite, which is smooth and less apt to form debris from the matrix. 8
ACCEPTED MANUSCRIPT Correspondingly, its counterface (See Fig. 9f) is well protected by the uniform and continuous transfer film due to the deposition of PTZ@RGO on the counterpart steel ring. Its tribological mechanism belongs to the typical abrasive wear [42]. These SEM images are consistent with the results of friction coefficient of the RGO/BMI and PTZ@RGO/BMI composites (See Fig. S4).
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To further illustrate the friction mechanism of PTZ@RGO/BMI and the advantages of PTZ@RGO in friction tests, we provide the illustration of working mechanism of RGO/BMI and PTZ@RGO/BMI (See Fig. 10). It can be seen in Fig. 10a and Fig. 10b that the RGO is easier to aggregate in BMI matrix than PTZ@RGO due to its poor compatibility with polymer matrix. During the wear process of the
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RGO/BMI composite, the aggregate RGO and the resin can be precipitated easily, then rough and discontinuous thin “carbon film” containing RGO and the BMI resin forms on the surface of counterpart
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steel ring, and which is difficult to effectively reduce the friction coefficient and wear rate of the RGO/BMI composite (See Fig. 10c). Compared to RGO/BMI composite, PTZ@RGO can be uniformly dispersed in the BMI matrix and form a good interfacial interaction with it (See Fig. 10d). Then more uniform and lubricating thinner “carbon film” containing PTZ@RGO forms on the surface of counterpart steel ring during the wear process of the composite, and effectively reduce the friction coefficient and
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wear rate of the PTZ@RGO/BMI composite. With the formation of the uniform and lubricating thinner “carbon film”, sliding occurs between the surface of the composite block and the “carbon film” other than between the rough surface of the composite and the steel counterpart [43].
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The outstanding tribological properties of the PTZ@RGO/BMI composites are mainly attributed to two aspects. Firstly, graphene itself has high strength and self-lubricating effect. It reinforces the composites, and their strengthened structure is bearing main body to increase abrasive resistance [44].
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Secondly, the active functional groups on the surface make graphene and BMI matrix have good compatibility and interfacial interaction, which enhances the interface strength. Load is easily to be transferred between polymer and graphene, thereby friction property is enhanced. Therefore, it can be concluded that PTZ@RGO can significantly improve the tribological performance of the BMI composites [45]. Thermal stability is of great importance for polymeric materials during the wear process. In present study, the thermal stability of neat BMI resin and its composites with 0.6 wt% RGO and PTZ@RGO was investigated by TGA using heating rate of 10 oC / min in nitrogen. As shown in Fig. 11a, all the samples 9
ACCEPTED MANUSCRIPT display similar degradation profile within the experimental temperature range, suggesting that the existence of fillers does not significantly alter the degradation mechanism of the BMI matrix. Fig. 11b, 11c and 11d show the maximum degradation temperature (Tm) of the BMI composites. The Tm values of the neat BMI, 0.6 wt% RGO/BMI and 0.6 wt% PTZ@RGO/BMI composites are 436.7, 448.6 and
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455.9oC, respectively. The improvement of the thermal stability of graphene-BMI composites may have originated from the outstanding barrier properties of graphene. Various studies have demonstrated that graphene with high aspect ratio and surface area in polymer provides a tortuous path for the diffusion of gas molecules and significantly reduces the permeation rate of gas [46, 47]. So, the graphene nanosheets
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act as physical barrier and delay the escape of degradation products. The interface interaction between the PTZ@RGO and BMI matrix can increase the thermal degradation activation energy of PTZ@RGO/BMI
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composites through restricting the thermal motion of polymer chains, thus improve the thermal stability of BMI resin. The excellent thermal stability can protect the composites from the harm of high temperature during the wear process, thus adhesive wear can be inhibited effectively [48, 49]. A large amount of heat can be generated during the friction process, which has a major impact on the physical and chemical structure of the material. So the friction final temperature of BMI composites
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friction surface with different filler loadings was measured by digital thermometer. It can be seen in Fig. 12 that the friction final temperature of the BMI composites gradually decrease as the increasing of the amount of fillers. This is mainly due to the fact that the presence of graphene not only has good
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self-lubricating property, but also can effectively transfer the heat generated during friction process, so as to prevent the deterioration of material performance due to excessive temperature, and thus improving the
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friction and wear resistance of the samples as a whole [50, 51]. When the content of fillers is more than 0.2 wt％, the friction final temperature of the PTZ@RGO/BMI composites is lower than those of RGO/BMI. This may be attributed to that the uniform dispersion of PTZ@RGO in the BMI resin to form continuous heat conduction path, thereby effectively dissipating the heat generated during the friction process. Hardness is a measure of the material resistance to various kinds of permanent shape change when a force is applied and the hardness of material is closely related with tribological properties. In this study, the dependency of the hardness of the RGO/BMI and PTZ@RGO/BMI composites on the content of the fillers is shown in Fig. 13. It can be observed that similar variation tendency in the hardness of the 10
ACCEPTED MANUSCRIPT RGO/BMI and PTZ@RGO/BMI-BA composites when the content is lower than 0.6 wt%. However, when the content of fillers is further increased, the hardness of the composite decreased, but is still higher than those of the neat BMI resin. This is attributed to the increasing of the addition amount, more graphene aggregation in the resin, resulting in more small voids in the material, which decreases the hardness [52,
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53]. 4. Conclusion
In this contribution, the reduced graphene oxide (RGO), cyanuric chloride (CC) and hexamethylenediamine (HMD) have been used to prepare the polytriazine@RGO nanocomposite
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(PTZ@RGO) with terminal amine by the method of one-step precipitation polymerization. The PTZ@RGO hybrid nanocomposite was employed as a kind of novel filler to improve the mechanical and
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tribological properties of BMI composite. The mechanical and tribological properties as well as thermal properties of the BMI composites can be improved by a very low loading of RGO or PTZ@RGO. Moreover, the PTZ@RGO/BMI composites have much better mechanical and tribological properties than those of RGO filler composites at the same filler content. This is for the reason that the PTZ on the graphene nanosheets surface influences the dispersibility of graphene in the polymer matrix as well as the
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interface interaction between the hybrid filler and the polymer matrix. This work has broad application potential in graphene-based nanocomposites where intimate interfacial bonding between the host and the structural modifier is critical, especially for improving mechanical, tribological and thermal properties of
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polymer matrices. Acknowledgement
This work was supported by the Natural Science Basic Research Plan in Shaanxi Province of China
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(Program No. 2018JQ5211), National Key Research and Development Program of China (No. 2017YFB0308602) and the Special Scientific Research Program Founded by Shaanxi Provincial Education Department (No. 18JK0102). References
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ACCEPTED MANUSCRIPT Figure Captions Fig. 1. Schematic diagram of the preparation process of PTZ@RGO and PTZ@RGO/BMI composites.
Fig. 3. XRD patterns of GO, RGO and PTZ@RGO.
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Fig. 2. FT-IR spectra of GO and PTZ@RGO.
Fig. 4. (a) XPS spectra of GO and PTZ@RGO; (b) C1s of GO; (c) C1s of PTZ@RGO. Fig. 5. TEM images of GO (a), RGO (b) and PTZ@RGO (c), respectively.
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Fig. 6. Relationship between mechanical properties and fillers content: flexural strength (a) and impact strength (b) of the composites, respectively.
Fig. 7. SEM images of fracture surfaces of neat BMI (a), RGO (0.6 wt%)/BMI (b), PTZ@RGO (0.6
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wt%)/BMI (c) and PTZ@RGO (1.0 wt%)/BMI (d) composites, respectively.
Fig. 8. Variation in mean coefficient of friction (a) and volume wear rate (b) of the composites: load, 198 N; sliding speed, 200 rpm/min; test duration, 120 min.
Fig. 9. The worn surfaces and counterface SEM images of neat BMI (a) and (d), RGO/BMI (b) and (e), PTZ@RGO/BMI (c) and (f) with 0.6 wt% filler content, respectively.
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Fig. 10. The friction mechanism simulation diagrams of RGO/BMI (10a and 10c) and PTZ@RGO/BMI (10b and 10d) composites before and after friction test, respectively. Fig. 11. TGA curves of neat BMI and its composites (a); TGA and DTG curves of neat BMI (b), 0.6wt％
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RGO/BMI (c) and 0.6wt％ PTZ@RGO/BMI, respectively. Fig. 12. Friction final temperature of the BMI composites friction surface with different filler loadings.
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Fig. 13. Hardness of the composites containing different contents of fillers.
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Fig. 1. Schematic diagram of the preparation process of PTZ@RGO and
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PTZ@RGO/BMI composites.
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Fig. 2. FT-IR spectra of GO and PTZ@RGO.
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Fig. 3. XRD patterns of GO, RGO and PTZ@RGO.
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Fig. 4. (a) XPS spectra of GO and PTZ@RGO; (b) C1s of GO; (c) C1s of PTZ@RGO.
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Fig. 5. TEM images of GO (a), RGO (b) and PTZ@RGO (c), respectively.
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and impact strength (b) of the composites, respectively.
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Fig. 6. Relationship between mechanical properties and fillers content: flexural strength (a)
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Fig. 7. SEM images of fracture surfaces of neat BMI (a), RGO (0.6 wt%)/BMI (b),
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Fig. 8. Variation in mean coefficient of friction (a) and volume wear rate (b) of the
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Fig. 9. The worn surfaces and counterface SEM images of neat BMI (a) and (d), RGO/BMI
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Fig. 10. The friction mechanism simulation diagrams of RGO/BMI (10a and 10c) and PTZ@RGO/BMI (10b and 10d) composites before and after friction test, respectively.
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Fig. 11. TGA curves of neat BMI and its composites (a); TGA and DTG curves of neat BMI
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Fig. 12. Friction final temperature of the BMI composites friction surface with different
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Fig. 13. Hardness of the composites containing different contents of fillers.
S1359-8368(18)33035-X
DOI:
https://doi.org/10.1016/j.compositesb.2019.02.005
Reference:
JCOMB 6591
To appear in:
Composites Part B
Received Date: 14 September 2018 Revised Date:
25 December 2018
Accepted Date: 5 February 2019
Please cite this article as: Liu C, Dong Y, Lin Y, Yan H, Zhang W, Bao Y, Ma J, Enhanced mechanical and tribological properties of graphene/bismaleimide composites by using reduced graphene oxide with non-covalent functionalization, Composites Part B (2019), doi: https://doi.org/10.1016/ j.compositesb.2019.02.005. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Graphical Abstract (Figure)
ACCEPTED MANUSCRIPT Enhanced mechanical and tribological properties of graphene/bismaleimide composites by using reduced graphene oxide with non-covalent functionalization† Chao Liu a∗, Yufei Dong b, Yang Lin a, Hongxia Yan c, Wenbo Zhang a, Yan Bao b, Jianzhong Ma b
University of Science &Technology, Xi’an 710021, China
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a Shaanxi Collaborative Innovation Center of Industrial Auxiliary Chemistry and Technology, Shaanxi
b College of Bioresources Chemical and Materials Engineering, Shaanxi University of Science & Technology, Xi’an 710021, PR China
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c Key Laboratory of Space Applied Physics and Chemistry, Ministry of Education, Department of Applied Chemistry, School of Science, Northwestern Polytechnical University, Xi’an 710129, China ABSTRACT: In order to simultaneously improve the dispersibility of graphene in the bismaleimide
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(BMI) resin and the interfacial interaction between them, the polytriazine@RGO nanocomposite (PTZ@RGO) with terminal amine was obtained by the method of one-step precipitation polymerization. The PTZ@RGO hybrid nanocomposite was then chosen as filler to improve the mechanical and tribological properties of BMI resin. Through effective surface modification, the compatibility and interface strength between graphene and BMI matrix were significantly improved, leading to the flexural
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strength, impact strength, anti-friction, wear resistance and thermal stability of PTZ@RGO/BMI composites were significantly improved as compared to those of neat BMI resin and RGO/BMI composites. The other properties including hardness and thermal properties of composites were all
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improved with increasing amount of PTZ@RGO fillers. These high performance BMI composites have a great potential for application in aerospace, mechanical engineering and other automobile brake devices.
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Key words: Polymer-matrix composites (PMCs); Theromosetting resin; Mechanical properties; Wear.
1. Introduction
Bismaleimide (BMI) is a high-performance thermosetting polymer with outstanding thermal stability, high specific modulus, and upstanding mechanical properties under the condition of high temperature [1-4]. Meanwhile, it has the similar formability to that of epoxy resin. While, the friction coefficient and wear rate of it are higher when it is used as friction material. Which are the major causes for material and ∗
Corresponding author. E-mail address: [email protected] (Chao Liu) 1
ACCEPTED MANUSCRIPT energy losses in mechanical processes. In order to improve mechanical durability and energy efficiency, BMI needs to be modified to enhance its performance, especially the mechanical and tribological comprehensive properties [5-7]. Recently, in order to overcome those shortcomings, varieties lubrication fillers such as carbon
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nanotube, MoS2, graphene or their derivative were introduced into the BMI matrix by researchers [8-10]. Graphene, the 2-dimensional carbon allotrope with outstanding self-lubricating and physicochemical properties, has attracted great interest as self-lubricating solid or as additive for polymer based friction materials [11-13]. The introduction of graphene into polymer matrix can not only improve its tribological
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properties but also improve its mechanical properties [14-16]. However, neat graphene contains no surface functional groups and has very limited dispersibility in polymer matrix, seriously limiting its
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potential application in preparation of functional nanocomposites. The surface modification of graphene by grafting of polymers, and the applications of such polymer-grafted nanoparticles have become an area of major interest lately. We have reported various hyperbranched polymer chemically surface modified graphene, which is of great value for improving compatibility with polymer matrix, resulting in the excellent distribution and enhanced interfacial bonding between graphene and polymer matrix [17, 18].
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However, this method has many steps and the yield of the product is low. Meanwhile the modification of covalent bond can result in the destruction of graphene structure, thus reducing its performance. Recently, the method of fabricating the functionalized graphene by deposition of nanoparticles on graphene to
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prevent the overlay of graphene sheets and enhance its tribological properties has attracted increasing attention [19-21]. Nevertheless, an appropriate interface interaction is essential for effective interface load transfer and wear resistance reinforcements. Furthermore, the mobility of both graphene is likewise
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crucial for achieving toughening and lubricating properties of composites. To resolve these contradictions, one expects to design a rational interface structure to optimize the mechanical and tribological properties of graphene-polymer composites [22, 23]. In this contribution, the reduced graphene oxide (RGO), cyanuric chloride (CC) and hexamethylenediamine (HMD) are used to prepare the polytriazine@RGO nanocomposite (PTZ@RGO) with terminal amine by the method of one-step precipitation polymerization. The PTZ@RGO was incorporated into the BMI matrix to form PTZ@RGO/BMI binary composite (See Fig. 1). This method can not only effectively resolve the aggregation of graphene but also improve the compatibility and 2
ACCEPTED MANUSCRIPT interfacial adhesion between graphene and the polymer matrix. The effects of RGO and PTZ@RGO content on the mechanical and tribological properties of BMI matrix were investigated. The wear mechanical of the composites was discussed. 2. Experimental
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2.1 Reagents and materials Natural graphite flakes (325 mesh) was obtained from Qingdao Hensen Graphite Co., Ltd. The graphene oxide (GO) nanosheets were produced from natural graphite flakes by the modified Hummers method [24]. Cyanuric chloride (CC, 99%), hexamethylene-diamine (HMD), NaOH and triethylamine
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(TEA) were purchased from Aladdin Chemistry Co., Ltd. N, N-Dimethylformamide (DMF), hydrazine hydrate and anhydrous alcohol were purchased from Tianjin Tianda Chemical Co. Ltd. The 4,
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4`-Bismaleimidodiphenylmathane (BDM) was provided by Rongchang Ning research group at Northwestern Polytechnical University. Diallyl Bisphenol A (DBA) was purchased from Sigma-Aldrich. All reagents were of analytical grade and used as received without further purification. 2.2 Procedures
Typically, graphene oxide (GO) was synthesized from natural graphite flakes by a modified Hummers
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method [24]. For synthesizing PTZ@RGO nanocomposite: firstly, 0.2 g GO was dispersed in 200 mL deionized water and sonicated in an ultrasonic bath for 0.5 h. Then, 2 mL hydrazine hydrate was added into the above solution and the mixture was stirred constantly at 98 oC kept for 6 h to prepare the RGO.
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Secondly, the RGO (0.2 g), CC (0.8 g, 0.5 mol) and a small amount of NaOH were added into a 250 mL three-neck round bottom flask. A mixture of DMF (150 mL) and 2.0 mL of TEA as an acid-acceptor were subsequently added. After ultrasonic treatment (150 W) for 30 min, HMD (1.6 g, 1.7 mol) was then added.
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The solution was maintained at 85 °C for 3 h with the protection of nitrogen gas by ultrasonic treatment (150 W). The product was filtered and washed by ethanol several times, and dried at 60 °C in a vacuum overnight.
The BMI-based composites were prepared by mixing the pre-weighed quantities of nanofillers,
Diallyl Bisphenol A (DBA) and BDM (DBA and BDM with a mass ratio of 3:4). Then the mixture was heated to 135 °C till the mixture totally melting and the nanofillers dispersed uniformly. Finally, the mixture was cured following the schedule of 150 °C/2 h + 180 °C/2 h + 220 °C/4 h. A post curing process was 250 °C/4 h. The BMI resin-based materials were also prepared in similar procedures to perform 3
ACCEPTED MANUSCRIPT contrast experiments. 2.3 Characterization Flexural strength and impact strength were performed according to GB/T2567-2008. The samples dimension for impact test and flexural test were (80 ± 0.2) × (10 ± 0.2) × (4.0 ± 0.2) and (80 ± 0.2) × (15
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± 0.2) × (4.0 ± 0.2) mm3, respectively. Each value obtained represented the average of five samples. The friction and wear tests were conducted on a test machine (MM-200, load 196 N, speed 200 rpm/min) under dry condition (temperature: 20±5oC, humidity: 50±10%) according to the GBT/3960-1983 (Chinese Standard). Where the pin and friction pair are the composites specimen and
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#45 steel disk, respectively. Before each test, the counterpart steel ring and the materials were sanded using No. 600 water-abrasive paper. Then the steel ring and samples were cleaned by acetone. The
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calculation of volume wear rate can be detailed in supporting information 1.
The Fourier transform infrared (FT-IR) spectra of samples was carried obtained with a Bruker IFS66/S FT-IR spectrometer. XRD patterns of the powered were measured on a Philips X’Pert MPD system with Cu-Ka radiation. X-ray photoelectron spectra (XPS) were investigated using a PHI Quantum 2000 Scanning ESCA Microprobe system. All XPS spectra were corrected using the C1s line at 284.6 eV.
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Raman spectra were obtained carried out on a Horiba Scientific XploRa Raman spectrometer. The microstructures of the composite specimens were characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM) via a Hitachi S-4800 and a FEI Tecnai G2 F20 microscopy,
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respectively. The hardness of the composite specimens was performed on a MH-5 Vector microhardness meter at a load of 50 N and a loading time of 20 s. Thermogravimetric analysis (TGA) was conducted in TGA Q50 at a heating rate of 10 oC/min, in a nitrogen atmosphere. Friction final temperature of the
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composites friction surface was recorded by a TMP-2 multichannel temperature tester. 3. Results and Discussion
3.1 Characterization of PTZ@RGO nanocomposite The FT-IR spectra of GO and PTZ@RGO are shown in Fig. 2. For the GO, the absorption peaks of GO appears at 1056, 1638 and 3430 cm-1, which are attributed to the stretching vibration of C=C, C-O and -OH groups (or the physical adsorption of water) on GO surface, respectively [25]. In the spectrum of PTZ@RGO, the peaks at 1443 and 1551 cm-1 are assigned to the skeleton vibration of the triazine ring, respectively. Additionally, the peak at 1330cm-1 is due to the C-N stretching vibration of the carbon atoms 4
ACCEPTED MANUSCRIPT of triazine rings linked with the -NH-. However, there is only one absorption peak at 3497 cm-1 between 3400 and 3500 cm-1, which may be attributed to the overlapping of the absorption peak of water and -NH2. The series of the peaks suggest the existence of PTZ in the PTZ@RGO nanocomposite [26]. XRD analysis was performed to investigate the structures of the samples. Fig. 3 shows the XRD
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patterns of GO, RGO and PTZ@RGO. As shown in Fig. 3, the XRD spectrum of the GO shows one strong (001) characteristic peak at 11.3° with d-spacing of 0.78 nm. After being reduced, the d-spacing of GO is reduced from 0.78 to 0.36 nm, which confirms that most oxygenic groups are removed. The XRD spectrum of RGO shows a very broad diffraction peak at 24.3° and the diffraction peak at 11.3°
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disappears, confirming the formation of RGO [27]. Compared with the XRD pattern of RGO, no obvious diffraction peaks for PTZ@RGO have been observed, which may be due to the formation of the
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PTZ@RGO nanocomposite so that individual RGO remains isolated [28].
XPS characterization was employed to analyze the surface chemical compositions of RGO and PTZ@RGO. In Fig. 4a, the peaks at 285.1, 399.6 and 532.0 eV are attributed to Cl2p, C1s, N1s and O1s, respectively [29]. From the C1s spectrum shown in Fig. 4b, four peaks attributed to C=C at about 284.5 eV, C-OH at about 286.6 eV, C-O-C at about 287.5 eV, and C=O at about 288.7 eV were observed. As
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presented in Fig. 4c, the C1s XPS peak-fitting of PTZ@RGO shows six different peaks, centered at 284.5, 285.6, 287.5, 288.1 and 288.7 eV, corresponding to C-C, C-N, C-O-C, C=N and C=O groups, respectively. However, the peaks at 287.5 and 288.7 eV have almost disappeared, which confirms a considerable
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degree of reduction. The peaks at 285.6 and 288.1 eV confirm that the existence of PTZ [30]. These results are in good agreement with the results of FT-IR. In order to further investigate the morphologies of the prepared hybrids, HR-TEM images were
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recorded. As shown in Fig. 5a and 5b, the GO and RGO have a typical shape resembling the wrinkled thin paper flake, and each nanosheet is up to several microns in size. After depositing PTZ onto the surface of RGO, a significant change can be observed (See Fig. 5c). The image of PTZ@RGO nanosheets show a gray texture of organic materials uniformly covering the entire surface of graphene nanosheets. This is attributed to that the polytriazine framework is rich in conjugated structure, which makes it can be firmly anchored to the surface of graphene [31]. These TEM observations are consistent with the other results from the other characterization methods, again demonstrating the successful deposition of PTZ to the graphene surface. The PTZ@RGO also shows excellent thermal stability (See Fig. S1). 5
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clearly seen that the composites containing the PTZ@RGO exhibit higher strength values than those of the RGO at almost all the filler content (See Fig. 6a). The flexural strength of RGO/BMI increases continuously with the addition of RGO and reaches a maximum value of 146.8 MPa at 0.4 wt%. For 0.4
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wt% PTZ@RGO content, the flexural strength of its composites increased to 167.5 MPa, which is increased as much as 26.5% in comparison with that of neat BMI resin (132.4 MPa). This phenomenon is attributed to the unique sheets structure of hybrid graphene and the interfacial adhesion between
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PTZ@RGO and BMI matrix, which can greatly improve the flexural strength of BMI resin [31]. Fig. 6b shows the effect of the loading concentration of fillers on the impact strength of BMI composites. It shows that the suitable amount of fillers can properly improve the impact strength of BMI resin. The impact strength of the BMI composites respectively incorporating 0.6 wt% RGO and PTZ@RGO reaches the maximum value of 14.84 and 17.36 kJ/m2, increasing as much as 25.0 % and
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46.3 % compared to that of neat BMI resin (11.87 kJ/m2). Unfortunately, the flexural and impact strength of the composites decline when the amount of fillers is further increased, which may be due to the fact that excessive fillers can not be well dispersed in the BMI matrix and agglomerate to cluster [32].
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Consequently, the advantages of the fillers can not be fully utilized and the mechanical properties of the composites decrease with the uneven distribution of the fillers in the BMI matrix [33].
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To further research the fracture characteristics of composites, SEM images of the facture surfaces of neat BMI and its composites after impact testing are shown in Fig. 7. The crack propagation regions of neat BMI resin and its composites are different. The fracture surface of neat BMI resin showed in Fig. 7a is comparatively slippery, which exhibits a typical brittle feature. After the introduction of fillers, significant difference on the facture surfaces of composites is shown in Fig. 7b, 7c and 7d, respectively. The RGO is obviously protruded and aggregated on the fracture surface of 0.6 wt% RGO/BMI composite, which is indicative of poor compatibility and weak interfacial interaction between the RGO and BMI matrix (See Fig. 7b). As described in Fig. 7c, the fracture surface of 0.6 wt% PTZ@RGO/BMI composite is indented and containing many ductile sunken areas, indicating a typical tough fracture features, which 6
ACCEPTED MANUSCRIPT is considered from a result of the uniform distribution of PTZ@RGO in the BMI matrix and good interface interaction between PTZ@RGO and BMI matrix. The results are consistent with those obtained by TEM (See Fig. S2). However, when the content of PTZ@RGO is further increased, the fillers agglomerated in the matrix (See Fig. 7d), and induced a deterioration of the mechanical properties.
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There are two main differences between RGO and PTZ@RGO, which create the different mechanical properties of the composites. On the one hand, the existence of the abundant polar functional groups on the PTZ@RGO surface gives rise to a homogenous distribution of PTZ@RGO in BMI matrix. This state of dispersion probably results in enhancement of the mechanical interlocking between
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PTZ@RGO and BMI polymer chains [34, 35]. These strong interfacial interactions enable graphene to bring its amazing mechanical properties into resin matrix, achieving a perfect reinforcing and toughening
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material [36]. On the other hand, compared with the relatively inert surface chemistry of RGO, PTZ@RGO owns abundant amino groups which can form covalent bonds with the carbon-carbon double bonds of BMI (The reaction mechanism of PTZ and BMI is shown in Fig. S3). The corresponding mechanism of the reaction has been discussed in our previous studies [37]. These chemical bonds contribute to the efficient load transfer between matrix and fillers, and play an excellent role in inhibiting
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crack propagation, thus resulting in higher impact strength [38]. And the good interface interaction between PTZ@RGO and BMI matrix can enhance the wear-resisting ability of the composites. Fig. 8a shows the friction coefficients of composites as a function of fillers content for steady-state
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sliding against the counterpart steel ring under dry conditions. It can be seen that, the friction coefficient of composites decreases significantly with the addition of fillers. When the content of RGO is 0.6 wt%, the friction coefficients of RGO/BMI composites reach the minimum value (0.26).The friction coefficient
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of PTZ@RGO/BMI composites reach the minimum value of 0.20 at 0.6 wt% PTZ@RGO, decreasing as much as 42.9% in comparison with that of neat BMI resin (0.35). This may be due to the fact that the fillers in the resin matrix near the worn surface is exposed and acted as the lubricating thin “carbon film” on the worn surface of the steel counterpart during the wear process. As the content of fillers exceeds 0.6 wt%, the friction coefficient of composites increases. But the friction coefficient of the composites is still lower than that of the neat BMI resin. This phenomenon can be explained that the excessive fillers agglomerates in the resin matrix and thus leads to more micro-cracks on the composites surface, and therefore reduces the surface smoothness of composites during the wear process [39]. At the same time, 7
ACCEPTED MANUSCRIPT the friction coefficient of PTZ@RGO/BMI composites is more stable than those of RGO/BMI composites
in the process of friction (See Fig. S4). This may be attributed to that the modification can improve the dispersibility of graphene in BMI resin, thus uniform self-lubricating transfer film can be formed on the surface of counterpart steel ring during the process of friction.
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Fig. 8b shows the volume wear rate of composites with different fillers content. It can be seen that the volume wear rate of composites decreases initially and then increased with the increasing of fillers. When the content of RGO is 0.6 wt%, the volume wear rate of the RGO/BMI composites reaches the lowest value by 4.12×10-6 mm3/ (N•m). The volume wear rate of PTZ@RGO/BMI composites also
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reaches the lowest value (1.72×10-6 mm3/ (N•m)) at the same fillers content, decreases as much as 89.2 % compared to that of the neat BMI resin (15.94×10-6 mm3/ (N•m)). This result is in agreement with the test
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results of impact strength, which mainly attributed to composites containing 0.6wt％ fillers can absorb more energy and increase the load-carrying capability of the composites, thus the deformation of the composites could be reduced effectively during the wear process [40]. However, the volume wear rates of composites slightly increase when the fillers content exceeds 0.6 wt%. This may be due to that excessive
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fillers agglomerate in the BMI matrix, thus reduces the load-carrying capability of the composites during the wear process [41]. The tribological properties of 0.6 wt% PTZ@RGO/BMI composite under different loads (See Fig. S5) and sliding speed (See Fig. S6) are also investigate. In order to further investigate the wear mechanism of the composites under dry friction condition,
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the worn surfaces morphologies of different types of composites and corresponding surfaces of counterpart steel ring were examined by SEM, respectively. For neat BMI (See Fig. 9a), the worn surface
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of it is rough and some loose plate-like debris can be observed. On the surface of counterpart steel ring (See Fig. 9d), the transfer film is thick and discontinuous and piles of furrow can be seen. These phenomena reflect that the neat BMI has relatively poor wear resistance in its sliding against the steel counterpart and its tribological mechanism mainly follows adhesive wear and fatigue wear mechanism. As shown in Fig. 9b, the worn surface of 0.6 wt% RGO/BMI composite is compact and less apt to form debris from the matrix. The uneven areas on the worn surface of RGO/BMI composites and the discontinuous transfer film on the surface of counterpart steel ring (See Fig. 9e) may be attributed to the non-uniform distribution of RGO in the BMI matrix. Fig. 9c shows the worn surface of 0.6 wt% PTZ@RGO/BMI composite, which is smooth and less apt to form debris from the matrix. 8
ACCEPTED MANUSCRIPT Correspondingly, its counterface (See Fig. 9f) is well protected by the uniform and continuous transfer film due to the deposition of PTZ@RGO on the counterpart steel ring. Its tribological mechanism belongs to the typical abrasive wear [42]. These SEM images are consistent with the results of friction coefficient of the RGO/BMI and PTZ@RGO/BMI composites (See Fig. S4).
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To further illustrate the friction mechanism of PTZ@RGO/BMI and the advantages of PTZ@RGO in friction tests, we provide the illustration of working mechanism of RGO/BMI and PTZ@RGO/BMI (See Fig. 10). It can be seen in Fig. 10a and Fig. 10b that the RGO is easier to aggregate in BMI matrix than PTZ@RGO due to its poor compatibility with polymer matrix. During the wear process of the
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RGO/BMI composite, the aggregate RGO and the resin can be precipitated easily, then rough and discontinuous thin “carbon film” containing RGO and the BMI resin forms on the surface of counterpart
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steel ring, and which is difficult to effectively reduce the friction coefficient and wear rate of the RGO/BMI composite (See Fig. 10c). Compared to RGO/BMI composite, PTZ@RGO can be uniformly dispersed in the BMI matrix and form a good interfacial interaction with it (See Fig. 10d). Then more uniform and lubricating thinner “carbon film” containing PTZ@RGO forms on the surface of counterpart steel ring during the wear process of the composite, and effectively reduce the friction coefficient and
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wear rate of the PTZ@RGO/BMI composite. With the formation of the uniform and lubricating thinner “carbon film”, sliding occurs between the surface of the composite block and the “carbon film” other than between the rough surface of the composite and the steel counterpart [43].
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The outstanding tribological properties of the PTZ@RGO/BMI composites are mainly attributed to two aspects. Firstly, graphene itself has high strength and self-lubricating effect. It reinforces the composites, and their strengthened structure is bearing main body to increase abrasive resistance [44].
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Secondly, the active functional groups on the surface make graphene and BMI matrix have good compatibility and interfacial interaction, which enhances the interface strength. Load is easily to be transferred between polymer and graphene, thereby friction property is enhanced. Therefore, it can be concluded that PTZ@RGO can significantly improve the tribological performance of the BMI composites [45]. Thermal stability is of great importance for polymeric materials during the wear process. In present study, the thermal stability of neat BMI resin and its composites with 0.6 wt% RGO and PTZ@RGO was investigated by TGA using heating rate of 10 oC / min in nitrogen. As shown in Fig. 11a, all the samples 9
ACCEPTED MANUSCRIPT display similar degradation profile within the experimental temperature range, suggesting that the existence of fillers does not significantly alter the degradation mechanism of the BMI matrix. Fig. 11b, 11c and 11d show the maximum degradation temperature (Tm) of the BMI composites. The Tm values of the neat BMI, 0.6 wt% RGO/BMI and 0.6 wt% PTZ@RGO/BMI composites are 436.7, 448.6 and
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455.9oC, respectively. The improvement of the thermal stability of graphene-BMI composites may have originated from the outstanding barrier properties of graphene. Various studies have demonstrated that graphene with high aspect ratio and surface area in polymer provides a tortuous path for the diffusion of gas molecules and significantly reduces the permeation rate of gas [46, 47]. So, the graphene nanosheets
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act as physical barrier and delay the escape of degradation products. The interface interaction between the PTZ@RGO and BMI matrix can increase the thermal degradation activation energy of PTZ@RGO/BMI
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composites through restricting the thermal motion of polymer chains, thus improve the thermal stability of BMI resin. The excellent thermal stability can protect the composites from the harm of high temperature during the wear process, thus adhesive wear can be inhibited effectively [48, 49]. A large amount of heat can be generated during the friction process, which has a major impact on the physical and chemical structure of the material. So the friction final temperature of BMI composites
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friction surface with different filler loadings was measured by digital thermometer. It can be seen in Fig. 12 that the friction final temperature of the BMI composites gradually decrease as the increasing of the amount of fillers. This is mainly due to the fact that the presence of graphene not only has good
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self-lubricating property, but also can effectively transfer the heat generated during friction process, so as to prevent the deterioration of material performance due to excessive temperature, and thus improving the
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friction and wear resistance of the samples as a whole [50, 51]. When the content of fillers is more than 0.2 wt％, the friction final temperature of the PTZ@RGO/BMI composites is lower than those of RGO/BMI. This may be attributed to that the uniform dispersion of PTZ@RGO in the BMI resin to form continuous heat conduction path, thereby effectively dissipating the heat generated during the friction process. Hardness is a measure of the material resistance to various kinds of permanent shape change when a force is applied and the hardness of material is closely related with tribological properties. In this study, the dependency of the hardness of the RGO/BMI and PTZ@RGO/BMI composites on the content of the fillers is shown in Fig. 13. It can be observed that similar variation tendency in the hardness of the 10
ACCEPTED MANUSCRIPT RGO/BMI and PTZ@RGO/BMI-BA composites when the content is lower than 0.6 wt%. However, when the content of fillers is further increased, the hardness of the composite decreased, but is still higher than those of the neat BMI resin. This is attributed to the increasing of the addition amount, more graphene aggregation in the resin, resulting in more small voids in the material, which decreases the hardness [52,
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53]. 4. Conclusion
In this contribution, the reduced graphene oxide (RGO), cyanuric chloride (CC) and hexamethylenediamine (HMD) have been used to prepare the polytriazine@RGO nanocomposite
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(PTZ@RGO) with terminal amine by the method of one-step precipitation polymerization. The PTZ@RGO hybrid nanocomposite was employed as a kind of novel filler to improve the mechanical and
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tribological properties of BMI composite. The mechanical and tribological properties as well as thermal properties of the BMI composites can be improved by a very low loading of RGO or PTZ@RGO. Moreover, the PTZ@RGO/BMI composites have much better mechanical and tribological properties than those of RGO filler composites at the same filler content. This is for the reason that the PTZ on the graphene nanosheets surface influences the dispersibility of graphene in the polymer matrix as well as the
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interface interaction between the hybrid filler and the polymer matrix. This work has broad application potential in graphene-based nanocomposites where intimate interfacial bonding between the host and the structural modifier is critical, especially for improving mechanical, tribological and thermal properties of
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polymer matrices. Acknowledgement
This work was supported by the Natural Science Basic Research Plan in Shaanxi Province of China
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(Program No. 2018JQ5211), National Key Research and Development Program of China (No. 2017YFB0308602) and the Special Scientific Research Program Founded by Shaanxi Provincial Education Department (No. 18JK0102). References
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ACCEPTED MANUSCRIPT Figure Captions Fig. 1. Schematic diagram of the preparation process of PTZ@RGO and PTZ@RGO/BMI composites.
Fig. 3. XRD patterns of GO, RGO and PTZ@RGO.
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Fig. 2. FT-IR spectra of GO and PTZ@RGO.
Fig. 4. (a) XPS spectra of GO and PTZ@RGO; (b) C1s of GO; (c) C1s of PTZ@RGO. Fig. 5. TEM images of GO (a), RGO (b) and PTZ@RGO (c), respectively.
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Fig. 6. Relationship between mechanical properties and fillers content: flexural strength (a) and impact strength (b) of the composites, respectively.
Fig. 7. SEM images of fracture surfaces of neat BMI (a), RGO (0.6 wt%)/BMI (b), PTZ@RGO (0.6
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wt%)/BMI (c) and PTZ@RGO (1.0 wt%)/BMI (d) composites, respectively.
Fig. 8. Variation in mean coefficient of friction (a) and volume wear rate (b) of the composites: load, 198 N; sliding speed, 200 rpm/min; test duration, 120 min.
Fig. 9. The worn surfaces and counterface SEM images of neat BMI (a) and (d), RGO/BMI (b) and (e), PTZ@RGO/BMI (c) and (f) with 0.6 wt% filler content, respectively.
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Fig. 10. The friction mechanism simulation diagrams of RGO/BMI (10a and 10c) and PTZ@RGO/BMI (10b and 10d) composites before and after friction test, respectively. Fig. 11. TGA curves of neat BMI and its composites (a); TGA and DTG curves of neat BMI (b), 0.6wt％
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RGO/BMI (c) and 0.6wt％ PTZ@RGO/BMI, respectively. Fig. 12. Friction final temperature of the BMI composites friction surface with different filler loadings.
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Fig. 13. Hardness of the composites containing different contents of fillers.
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Fig. 1
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Fig. 1. Schematic diagram of the preparation process of PTZ@RGO and
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Fig. 2. FT-IR spectra of GO and PTZ@RGO.
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Fig. 3. XRD patterns of GO, RGO and PTZ@RGO.
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Fig. 3
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Fig. 4
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Fig. 4. (a) XPS spectra of GO and PTZ@RGO; (b) C1s of GO; (c) C1s of PTZ@RGO.
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Fig. 5
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Fig. 5. TEM images of GO (a), RGO (b) and PTZ@RGO (c), respectively.
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Fig. 6
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and impact strength (b) of the composites, respectively.
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Fig. 6. Relationship between mechanical properties and fillers content: flexural strength (a)
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Fig. 7. SEM images of fracture surfaces of neat BMI (a), RGO (0.6 wt%)/BMI (b),
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Fig. 8. Variation in mean coefficient of friction (a) and volume wear rate (b) of the
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composites: load, 198 N; sliding speed, 200 rpm/min; test duration, 120 min.
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Fig. 9. The worn surfaces and counterface SEM images of neat BMI (a) and (d), RGO/BMI
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(b) and (e), PTZ@RGO/BMI (c) and (f) with 0.6 wt% filler content, respectively.
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Fig. 10. The friction mechanism simulation diagrams of RGO/BMI (10a and 10c) and PTZ@RGO/BMI (10b and 10d) composites before and after friction test, respectively.
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Fig. 11
Fig. 11. TGA curves of neat BMI and its composites (a); TGA and DTG curves of neat BMI
AC C
EP
TE D
(b), 0.6wt％ ％ RGO/BMI (c) and 0.6wt％ ％ PTZ@RGO/BMI, respectively.
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RI PT
Fig. 12
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Fig. 12. Friction final temperature of the BMI composites friction surface with different
AC C
EP
TE D
filler loadings.
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RI PT
Fig. 13
AC C
EP
TE D
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Fig. 13. Hardness of the composites containing different contents of fillers.