Effects of 2-D transition metal carbide Ti₂CT_x on properties of epoxy composites

Abstract

Here, 2-D nanostructured Ti₂CT_x was successfully synthesized by selectively exfoliating the Ti₂AlC MAX phase via a liquid etching technique. Investigated as a promising and functional reinforcement additive, 2-D Ti₂CT_x nanosheets were dispersed into an epoxy matrix evenly at different mass percentages to obtain self-lubricating and anti-friction Ti₂CT_x/epoxy resin (EP) nanocomposites with high toughness and low creep strain by in situ intercalative polymerization. The phase structure and morphology of Ti₂CT_x and Ti₂CT_x/EP nanocomposites were investigated by X-ray diffraction (XRD) and scanning electron microscopy (SEM). Subsequently, this work studied the effect of Ti₂CT_x on the EP matrix in terms of fracture surface morphology, glass transition temperature (T_g), storage modulus (E_s), thermal stability, mechanical properties and dry sliding anti-friction properties. The results show that Ti₂CT_x interlayers were intercalated and exfoliated with EP molecular chains and intertwined with each other, which produced a greater interfacial area between the inorganic additive and the polymer matrix, resulting in the formation of Ti₂CT_x/EP nanocomposites without adverse structural defects. With the addition of an appropriate mass fraction of Ti₂CT_x, the fracture toughness and flexural strength of EP composites reached 17.8 kJ m⁻² and 98 MPa, which represented increases of 76% and 66%, respectively, compared with those of pure EP. In addition, increases in T_g and E_s of 15% and 38% were achieved, respectively, relative to those of pure EP when the concentration of Ti₂CT_x was 2.0 wt%. Moreover, tribological analysis indicated that there was a great improvement in anti-friction properties and the morphology of worn surfaces of the nanocomposites appeared to be smoother than that of pure EP, which implied the transformation of fatigue and abrasion wear into adhesion wear. This work is of great significance in that novel 2-D transition metal carbide crystals were employed to fill and modify a thermoset polymer resin and enhance its toughness and strength and, as a self-lubricating additive, enabled a great breakthrough in the anti-friction properties of polymer-matrix composites.

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1. Introduction

Epoxy resin (EP), as a high-performance thermoset synthetic polymer resin, possesses prominent advantages such as a low shrinkage rate, high adhesive strength, corrosion resistance, strong dielectric properties, environmental stability, relatively low cost, etc., and has been used in the design of intricate structural parts employed in aerospace and automobile components and electrical insulation.^1,2 In actual applications, higher requirements have been proposed for polymer materials that took the place of wood and steel. However, its high brittleness, low toughness and resistance to thermal shock, and poor friction and wear properties have restricted the development of EP in applications, especially in anti-friction and high-frequency vibrating parts in machinery. Therefore, it is a difficult problem to increase the toughness and enhance the creep resistance and anti-friction properties of EP by physical or chemical methods.

Currently, the improvement of polymer performance via adding fillers to synthesize polymer-matrix composites has become a research focus.³ Taking into consideration the excellent properties of inorganic nanomaterials, they are widely used as fillers to improve the properties of polymers. Mayoral et al. synthesized graphene nanoplatelet (GNP)/polyamide 6 composites by melt processing, which displayed a 420% increase in Young's modulus at a 20 wt% loading of nanofiller.⁴ Chemical oxidative polymerization has been employed to synthesize magnetite/SPAN conductive polymer-based nanocomposites, which have satisfactory conductivity.^5,6 Choi et al. employed a sol–gel method to prepare graphene (FGS)/waterborne polyurethane nanocomposites, in which the finer dispersion of FGS improved the conductivity and thermal resistance of polyurethane and acted as a reinforcing filler to increase the tensile strength and modulus.⁷ Ultrasonic irradiation polymerization was employed to synthesize nano-SiO₂/poly(3-aminophenylboronic acid) (PAPBA) composites with a core–shell morphology, which exhibited a high conductivity of 0.2 S cm⁻¹.⁸ In addition, microemulsion polymerization is also a vital way of preparing organic–inorganic nanocomposites. Reddy et al.⁹ reported that microemulsion polymerization was employed to coat MWCNTs with polyaniline (PANI) nanospheres and proved that a MWCNTs/PANI nanocomposite without a surfactant had higher conductivity than a composite synthesized by in situ chemical polymerization. Besides, utilizing in situ chemical oxidative graft polymerization,¹⁰ a hybrid MWCNTs-f-PANI-noble metal nanocomposite was synthesized and displayed significantly higher conductivity than those of CNTs-f-PANI and pure PANI.

However, the synthesis methods of epoxy-based nanocomposites have mainly employed direct mixing and intercalative polymerization because of the fluidity of epoxy monomers. Direct mixing is one of the easiest methods of synthesizing nanoparticles/EP composites¹¹ via sonication and mechanical stirring, which disperses surface-treated nanoparticles into an EP matrix and minimizes agglomeration. Owing to the unique microscopic morphology of two-dimensional (2-D) nanomaterials (e.g., graphene,^12–14 nanoclay,¹⁵ ceramics,² MoS₂,^16,17 BN,¹⁸ etc.), they can be exfoliated into monolithic nanolayers and filled into an EP matrix to manufacture high-performance composites by in situ intercalative polymerization. Via in situ intercalative polymerization, a 2-D layer nanostructure bonded by van der Waals bonds is intercalated by epoxy molecular chains, which results in an increase in the interlayer distance and separation into strips after curing.^3,19–21 Moreover, the exfoliated nanosheets embedded in the epoxy matrix could produce a larger interfacial surface area between the inorganic additive and the polymer matrix. Moreover, because in situ intercalation achieves exfoliation, 2-D lamellar nanosheets play a better role in interfacial interactions with an epoxy matrix.^15,22

Recently, new types of 2-D transition metal carbides and carbonitrides called MXenes, which possess graphene-like 2-D morphology^23,24 with a general formula of M_n+1X_nT_x, where M is an early transition metal, X is C and/or N, n is 1–3, T is a surface terminating functionality (O, F and/or OH), and x represents the number or type of functional groups,^23,25 were fabricated by selective etching of the A element (an A-group element, mostly group IIIA or IVA) layers from MAX phases, which are a large family of ternary layered carbide and nitride ceramics,²⁶ and attracted much attention and triggered numerous applications.^27–31 Owing to their novel electrical properties,^32–34 excellent mechanical strength,³⁵ adsorption properties,³⁶ etc., MXenes have been widely used in research into new energy materials,^37–41 lubricant additives,^42,43 and polymer-reinforced materials.^44,45

To our knowledge, hardly any literature has reported MXene/EP composites. Here, taking advantage of the layered structure and surface terminating functionalities of MXenes, Ti₂CT_x/EP composites were synthesized by means of in situ intercalative polymerization, as shown in Fig. 1, and the effects of the 2-D transition metal carbide Ti₂CT_x on the properties of epoxy-based composites were investigated in detail. Moreover, MXene nanosheets played a role as a reinforcing framework to support crosslinked EP molecular chains, as graphene⁴⁶ or the 2-D MAX phase Carbide Derived Ceramic Nanostructure (CDCN)² played the role in an EP matrix of increasing the adhesion force and acting as an interlocking chain between the nanofillers and the EP matrix in previously reported papers. A Ti₂CT_x/EP nanocomposite was prepared successfully, which had better mechanical, thermal, creep resistance and anti-friction properties than those of pure EP. These results show that 2-D MXene nanosheets have significant potential as reinforcing additives in polymer-based composites.

Fig. 1 Schematic diagram of EP interlayered into Ti₂CT_x layers: (a) Ti₂CT_x was dispersed into acetone and mixed with EP, where the two phases separated without interacting; (b) following ultrasonic vibration and mixing, EP monomer was intercalated into Ti₂CT_x layers; (c) cured epoxy resin with extensive separation of Ti₂CT_x into strips.

2. Experimental

2.1 Materials

The MAX phase Ti₂AlC, which is a precursor of the 2-D nanomaterial Ti₂CT_x, was synthesized by pressless sintering in our previous report.⁴⁷ The pristine Ti₂CT_x that was utilized in this study was synthesized by selectively exfoliating A atoms from the MAX phase in a solution of lithium fluoride (LiF, 99%, Tianjin Fine Chemical Co., Ltd, China) in hydrochloric acid (HCl, AR grade). Epoxy resin (DGEBA, epoxy equivalent 184–195 g mol⁻¹, viscosity 10–16 Pa s) was purchased from Xingchen Synthetic Material Co., Ltd (China) and 2-ethyl-4-methylimidazole (EMI-2,4, 98%, C₆H₁₀N₂, MW 110.16 g mol⁻¹) hardener and antifoaming agents (AF) were supplied by Chinese Macklin Biochemical Co., Ltd. Acetone and ethyl alcohol (AR grade, 99%) were supplied by Hongyan Reagent Company of China and distilled water (DW) was lab-synthesized.

2.2 Preparation of 2-D Ti₂CT_x MXene

The novel 2-D nanostructured MXene Ti₂CT_x was synthesized by solution etching.⁴⁸ Initially, 5.0 g LiF was dispersed into 100 mL 6 M HCl solution in a PTFE container. After ultrasonic concussion for 10 min, 5.0 g Ti₂AlC was weighed and added to the above solution over 5 min and ultrasonic vibration was continued for 30 min. Afterwards, the container was transferred to a magnetic stirring apparatus and stirred for 48 hours at 40 °C. The resulting suspension was subjected to ultrasonic treatment for 10 min and then centrifuged at 5000 rpm for 10 min. The supernatant was decanted and the remaining sediment was washed with DW several times until the pH value of the supernatant was close to neutral. Finally, the sediment was washed with ethyl alcohol three times and dried in a vacuum at 80 °C for 12 h.

2.3 Fabrication of Ti₂CT_x/EP nanocomposite

Initially, different masses (0.5 g, 1.0 g, 1.5 g and 2.0 g, respectively) of the prepared Ti₂CT_x were weighed, dispersed into 10 mL acetone and subjected to ultrasonic vibration for 30 min at room temperature (RT). Then EP (100 g) and AF (0.25 g) were added to the above solutions with simultaneous mechanical stirring at a speed of 1500 rpm and ultrasonic treatment for 1 hour at RT and then heated to 50 °C with continued stirring at the same speed, again with ultrasonic treatment, for 30 min. An intermediate degassing procedure was carried out at 60 °C for 5 hours to remove acetone and trapped air. Afterwards, the hardener EMI-2,4 (5 g) was added and stirring and ultrasonic concussion were continued for one hour, then the mixture was degassed by vacuum-pumping for 2 hours at RT. The resulting mixture was highly degassed, followed by being poured into a PTFE mould, which had been preheated at 80 °C in order to ensure better wetting between the mixture and the mould. Afterwards, the curing process included three stages, namely, pre-curing at 60 °C for 1 h, curing at 100 °C for 2 h and post-curing at 120 °C for 2 h. The whole process was carried out in a high-temperature drying oven.

2.4 Characterization and measurement

2.4.1 Crystal phase and microstructure observation. X-ray diffraction (XRD, SmartLab X-ray diffractometer equipped with copper target Kα radiation, λ = 1.5406 Å, Rigaku Corporation, Japan) was used to detect the crystal structures of Ti₂AlC, Ti₂CT_x powder and Ti₂CT_x/EP nanocomposites. The samples were scanned over the two theta angle range (2θ) of 2–80° in steps of 0.02° and at a scanning speed of 10° per min to obtain the diffraction patterns.

The morphology and microstructures of the powders and nanocomposites were examined by a field emission gun scanning electron microscope (FEG-SEM, accelerating voltage 10 kV, FEI Quanta 250, Germany). The samples were fixed on a flat metal pallet with conductive adhesive tape and the test surface was sprayed with gold for 90 seconds.

2.4.2 Dynamic mechanical and thermal analysis. Dynamic mechanical analysis (DMA) was conducted using a dynamic mechanical analyzer (DMA Q800, TA Instruments, Inc., USA) equipped with a single cantilever clamp, and the rectangular test specimens had dimensions of 35 mm (length) × 12.5 mm (width) × 2.0 mm (thickness). The test was carried out by a temperature sweep (temperature ramp from 50 °C to 220 °C at 3° per min) method at a constant frequency of 1 Hz to obtain the dynamic storage modulus and damping curves (tan delta) of the specimens. The T_g values obtained from the maximum peaks in the tan delta curves correspond to the T_g temperature.

Creep and recovery tests were performed using the above-mentioned DMA with the same sample dimensions. The creep properties were characterized as a function of creep time for 20 min and recovery time for 20 min under applied stress levels of 5, 10 and 15 MPa at 75 °C.

Thermal gravimetric analysis (TGA) of samples was carried out using a Setaram Evolution 2400 thermal analyzer under an argon flow of 20 mL min⁻¹ with a heating rate of 5 °C min⁻¹ from room temperature to 600 °C.

2.4.3 Mechanical properties measurement. The microhardness of pure EP and Ti₂CT_x/EP nanocomposites was investigated using a Q/NFAK₁-2000 Vickers microhardness tester (MH-5, Shanghai Hengyi Electronic Testing Equipment Co., Ltd, China) with a Vickers diamond pyramidal indenter having a square base and a pyramidal angle of 136°. The specimens were subjected to a load of 2 N with a dwell time of 20 s and received indentations. The microhardness of each specimen was measured five times and the results were averaged.

The impact strength was measured by a Charpy impact strength tester (XJ-50Z, Combination of Impact Testing Machine, China), which was equipped with a pendulum with an impact energy of 7.5 J and a blade angle of 30°, which fell from an elevation of 160° with an impact velocity of 3.8 m s⁻¹ to fracture the specimens to obtain the value of the impact energy. The dimensions of the specimens were 50 mm (length) × 6.0 mm (width) × 4.0 mm (thickness). The results for the value of the impact energy were obtained by measuring three times and averaging.

Flexural strength tests on the composites were performed with a WDW-20 Electronic Universal Testing Machine (China) in three-point bending test mode, which had a load cell capacity of 10 kN. The elongated test specimens were prepared with dimensions of 40 mm (length) × 10.0 mm (width) × 2.0 mm (thickness). The span distance of the three-point bending test was 32 mm. The crosshead had a speed of 2 mm min⁻¹. The values of the flexural strength were obtained by measuring three times and then averaging.

2.4.4 Anti-friction measurements. The tribological properties of Ti₂CT_x/EP composites were investigated with pin-on-disk pairs provided by a multi-specimen test machine, which was consistent with our previous method reported in the literature.⁴⁵ As the disk pairs, specimens with dimensions of 40 mm × 40 mm × 4.0 mm and the frictional surfaces of steel pins and disks were polished with 1200 mesh SiC metallographic sandpaper and washed with acetone. Friction testing was conducted at rotating speeds of 100 rpm (linear speed 0.15 m s⁻¹) and 200 rpm (linear speed 0.3 m s⁻¹) under loads of 98 N for a test duration of 30 min to investigate the effect of the addition of different amounts of Ti₂CT_x and the slipping velocity on the frictional coefficient of Ti₂CT_x/EP nanocomposites.

3. Results and discussion

3.1 Polymerization mechanism for the formation of composites

The typical polymerization mechanism of DGEBA cured by EMI-2,4 hardener is shown in Fig. 2. Firstly, the hardener acting as a catalyst gave rise to ring-opening and breakage of C–O bonds in epoxy monomer and moreover N–H bonds are broken and H is transferred from N–H to O to form the intermediate. Secondly, the tertiary amine initiated the ring-opening of EP monomer and began the polymerization reaction. On the addition of Ti₂CT_x, we inferred that the surface OH functional groups bonded to the epoxy groups with simultaneous ring-opening, in a similar way to an etherification reaction. In other words, the addition of Ti₂CT_x played the role of a bridge that promoted the curing reaction and the formation of a crosslinked network in the EP matrix. As shown by the results of FT-IR spectroscopy in Fig. 3, the fingerprint peak of epoxy groups in epoxy monomer at 915 cm⁻¹ is significant; after curing, the disappearance of the fingerprint peak means that the initiator has broken C–O bonds in the epoxy group and induced ring-opening. Moreover, peaks due to OH group stretching vibrations appeared at 3600–3800 cm⁻¹ for the cured EP and Ti₂CT_x/EP nanocomposite,² corresponding to the ring-opening of epoxy groups and introduction of OH groups in the curing process in Fig. 2.

Fig. 2 Reaction mechanisms and polymerization process of DGEBA cured by EMI-2,4 hardener.

Fig. 3 FT-IR spectra of epoxy monomer before and after curing and Ti₂CT_x/EP composites with the addition of 2.0 wt% Ti₂CT_x.

3.2 Phase and morphology analysis of Ti₂CT_x and Ti₂CT_x/EP nanocomposites

Fig. 4 shows the XRD patterns of MAX phase Ti₂AlC, MXene Ti₂CT_x, pure EP and a Ti₂CT_x/EP composite. Obviously, we can see that the diffraction peaks of the (002), (103), and (106) crystal planes of Ti₂AlC, which corresponded to values of 2 theta of 12.9°, 39.4°, and 53.1°, respectively, disappeared after liquid etching. Simultaneously, a new diffraction peak appeared at 7.37°, which corresponded to the (002) crystal plane of Ti₂CT_x.⁴⁹ Moreover, the XRD pattern of pure EP indicated its amorphous nature based on the broad diffraction hump at a 2 theta angle of about 17.5°.⁵⁰ It can be seen in the inset of Fig. 4 that the d-spacing of Ti₂CT_x for (002) crystal planes increased from 11.7 Å to 14.4 Å and correspondingly the 2 theta angle shifted from 7.37° to 6.23°; in other words, the diffraction peak shifted left to a lower angle and revealed an increase in the layer spacing. In addition, the intensity of the (002) diffraction peak of Ti₂CT_x in the composite decreased and the full width at half maximum (FWHM) was significantly greater compared with that for Ti₂CT_x. The reason for this result is that chains of epoxy molecules were intercalated into the interlayer and interweaved with Ti₂CT_x, causing an increase in the layer spacing of Ti₂CT_x and even destruction of the crystal structure, as in the observation that the structures of silicate clay or graphite were intercalated by molecular chains of epoxy resin in previous reports.^50,51

Fig. 4 XRD patterns of MAX phase Ti₂AlC, MXene Ti₂CT_x, pure EP and EP nanocomposite with a Ti₂CT_x content of 2.0 wt%.

Fig. 5 shows a schematic diagram of the preparation process of Ti₂CT_x and Ti₂CT_x/EP composites. Fig. 5(a) explains the etching of the ternary layered ceramic MAX phase Ti₂AlC to give the MXene Ti₂CT_x at the level of atomic structure, in which the active metallic element Al in the MAX phase was exfoliated and the remaining atomic layers of transition metal and carbon were bonded by covalent bonds to form the stable 2-D transition metal carbide Ti₂CT_x, which is similar to treating the MAX phase with hydrofluoric acid.⁴⁹ Subsequently, EP molecular chains were intercalated into Ti₂CT_x layers and cured, resulting in an increase in the layer spacing. Fig. 5(b) and (c) shows SEM pictures of Ti₂AlC particles and 2-D layer-structure Ti₂CT_x, in which the thickness of a single layer is about 50 nm and the spacing between Ti₂CT_x layers is approximately 100 nm and molecular chains of epoxy resin were easily intercalated into the interlayers. Fig. 5(d) shows the fracture surface of the cured Ti₂CT_x/EP nanocomposite. It can be observed that layered Ti₂CT_x was exfoliated into a single layer or few layers and dispersed in the EP matrix homogeneously, which demonstrated the good compatibility of EP and Ti₂CT_x.

Fig. 5 (a) Schematic diagram of preparation of Ti₂CT_x and Ti₂CT_x/EP nanocomposite; SEM of MAX phase Ti₂AlC (b), which was etched to form 2-D layer-structure Ti₂CT_x (c) and a Ti₂CT_x/EP nanocomposite (d). The inset in (b), (c) and (d) is the high magnification image.

The fracture morphology of pure EP and Ti₂CT_x/EP nanocomposites is shown in Fig. 6. There only appear to be pores of a few microns in size on the relatively smooth fracture surface of pure EP (Fig. 6a), which exhibits brittle fracture characteristics. Fig. 6(b)–(d) shows the fracture surfaces of Ti₂CT_x/EP nanocomposites, which are rougher than that of pure EP owing to the addition of Ti₂CT_x. Clearly, Ti₂CT_x layers were intercalated and expanded (Fig. 6b), and were exfoliated into single layers and embedded in the EP matrix (Fig. 6c). Furthermore, it was demonstrated clearly that layered Ti₂CT_x was dispersed uniformly in the polymer matrix without agglomeration on a cross-section of the EP composite (Fig. 6c and d). In addition, the fracture surfaces of Ti₂CT_x/EP nanocomposites exhibited a typical flow pattern, which can be used to estimate the local direction of crack propagation, which is marked by the black arrows on the cross-sections of the composites.^52,53 With an increase in the Ti₂CT_x content, the number of micro-cracks increased and the fracture surface displayed typical tough fracture features. A similar phenomenon has been observed in previous research.^21,53,54

Fig. 6 FE-SEM images obtained of the fracture surfaces of cured samples: (a) pure EP, (b) molecular chains of EP intercalated into Ti₂CT_x layers resulting in an increase in the layer spacing, (c) and (d) cured epoxy resin with Ti₂CT_x sheets extensively exfoliated into single layers.

3.3 Analysis of thermal, mechanical and anti-friction properties of composites

3.3.1 Dynamic mechanical analysis and thermal stability. The thermal stability of polymer materials plays a vital role in engineering applications. In particular, the glass transition temperature (T_g) is one of the most critical parameters, which indicates the upper temperature limit for the application. Moreover, changes in T_g values have a crucial relationship with reinforcement phases and the interfacial interaction between reinforcement materials and the epoxy resin matrix.^15,55 Fig. 7(a) shows the loss factor (tan delta) of pure EP and EP composites with different mass fractions of Ti₂CT_x. The T_g value was determined from the temperature at the peak of the corresponding glass transition region of the tan delta curve, which increased gradually with a change in the mass concentration of Ti₂CT_x (Table 1). Obviously, the T_g of pure EP is 141.5 °C, whereas when the concentration of Ti₂CT_x increased to 2.0 wt%, the T_g of the Ti₂CT_x/EP nanocomposite increased by ∼20 °C and ∼15%, respectively. In other words, the addition of Ti₂CT_x can increase the T_g value dramatically, which can be explained by the interfacial interaction and crosslinking between Ti₂CT_x nanosheets and EP molecular chains. As mentioned above, EP molecular chains were grafted on the hydroxyl-functionalized surface of Ti₂CT_x nanosheets and in combination with the nanosheets interweaved into a three-dimensional crosslinked network structure during the curing process. When the temperature increased to the glass transition region, because the Ti₂CT_x surface developed a strong interaction with molecules in the matrix, stiff nanosheets hindered the vibration of molecular chains and caused an increase in the relaxation time, which led to an increase in the value of T_g.^15,56–58

Fig. 7 Results of DMA and TGA: (a) tan delta curve, from which was obtained the glass transition temperature (T_g) and (b) storage modulus (E_s), (c) TG and (d) DTG of pure EP and Ti₂CT_x/EP nanocomposites.

Table 1 Detailed parameters of the results of analysis of the thermal properties of Ti₂CT_x/EP nanocomposites with different amounts of Ti₂CT_x^a

Sample	DMA		TGA
	Tg (°C)	Es (GPa)	Tonset (°C)	Td10 (°C)	Tdm (°C)	ΔHd (J g^−1)
a Tg: glass transition temperature; Es: storage modulus at 50 °C; Tonset: starting degradation temperature; Td10: decomposition temperature of 10 wt% weight loss; Tdm: fastest decomposition temperature at the peak of the DTG curve; ΔHd: degradation enthalpy per gram.
Pure EP	141.5	1.33	383.6	401	438.2	46.9
Ti2CTx/EP (0.5%)	152.1	1.31	383.9	401	438.8	47.2
Ti2CTx/EP (1.0%)	155.0	1.82	373.4	399	437.6	40.0
Ti2CTx/EP (1.5%)	155.6	1.79	373.2	398	437.1	37.1
Ti2CTx/EP (2.0%)	162.1	1.77	372.7	395	435.9	35.6

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The storage modulus (E_s) reflects the viscoelasticity and rigidity of materials, which determine their ability to resist deformation when affected by external forces of constant frequency and temperature. Fig. 7(b) shows that as the temperature increased there was a change in curves of E_s versus different contents of Ti₂CT_x added to the EP matrix, indicating a marked increase in E_s (storage modulus at 50 °C; the values are listed in Table 1). Moreover, the viscoelastic transition temperature of EP composites increased with an increase in the content of Ti₂CT_x nanosheets. The main reason is probably that the reinforcing effect of Ti₂CT_x nanosheets and the large specific surface area and unique two-dimensional planar structure of Ti₂CT_x enhanced the adhesion and interlocking action between the nanofiller and the EP molecular chains and improved the interfacial interaction, resulting in a consequent increase in the toughness of Ti₂CT_x/EP composites (Table 2).¹²

Table 2 Detailed parameters of mechanical and tribological properties of Ti₂CT_x/EP nanocomposites with different amounts of Ti₂CT_x

Sample	Mechanical properties			Friction coefficient
	Microhardness (Hv)	Flexural strength (MPa)	Impact strength (kJ m^−2)	Sliding velocity
				0.15 (m s^−1)	0.3 (m s^−1)
Pure EP	23.8 ± 1.2	59 ± 1.4	10.1 ± 0.2	0.768	0.673
Ti2CTx/EP (0.5%)	19.6 ± 0.4	78 ± 1.7	15.2 ± 1.1	0.801	0.605
Ti2CTx/EP (1.0%)	18.2 ± 0.1	98 ± 2.0	17.8 ± 1.6	0.717	0.365
Ti2CTx/EP (1.5%)	18.6 ± 0.1	84.5 ± 2.4	14.6 ± 1.6	0.67	0.257
Ti2CTx/EP (2.0%)	20.8 ± 1.8	81 ± 1.1	13.3 ± 2.1	0.454	0.228

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Fig. 7(c) and (d) shows the TGA and DTG curves of EP composites with different contents of Ti₂CT_x, and the detailed parameters are listed in Table 1. Obviously, the starting degradation temperature (T_onset), fastest decomposition temperature (T_dm), 10% weight loss temperature (T_d10) and degradation enthalpy per gram (ΔH_d) generally decreased as the content of Ti₂CT_x in the composites increased. These results can be explained because Ti₂CT_x nanosheets with high thermal conductivity were finely dispersed in the EP matrix and displayed a rapid thermal response;⁵⁹ when the temperature rose the motion and rearrangement of EP molecular chains were hindered by the dispersed Ti₂CT_x nanosheets⁶⁰ owing to their mutual interactions. Therefore, the addition of Ti₂CT_x nanosheets will possibly reduce the thermal stability of the composites to a certain extent.

3.3.2 Creep and recovery behavior. As we know, it is extremely important to improve the creep resistance of polymer materials, in particular for their long-term use under a certain load and temperature. Fig. 8 shows the creep and recovery properties of pure EP and Ti₂CT_x/EP composites with different contents of Ti₂CT_x under different applied stresses in relatively high-temperature conditions. Clearly, a higher level of creep stress resulted in larger strain in pure EP and its composites, as shown in Fig. 8(a)–(c). Without regard to the change in levels of applied stress, compared with that of pure EP, the creep strain of the composites underwent a clear reduction with an increase in Ti₂CT_x content. Furthermore, it is notable that the instantaneous and highly elastic deformation of Ti₂CT_x/EP composites underwent a greater recovery after removal of the applied stress, and the composites exhibited a lower permanent deformation than pure EP, which indicated that there was a significant improvement in creep recovery performance with the addition of Ti₂CT_x nanosheets to the EP matrix. Fig. 8(d) shows the creep strain and unrecovered strain of pure EP and its composites with different contents of added Ti₂CT_x under high levels of applied stress at 75 °C. It is worth noting that the creep strain and unrecovered strain of the composite with 2.0 wt% Ti₂CT_x nanosheets after a creep time of 20 min at a stress level of 15 MPa and removal of stress with a recovery time of 20 min displayed reductions of ∼20% (creep strain ∼1.1%) and ∼40% (unrecovered strain ∼0.06%) compared with those of pure EP (∼1.33% and ∼0.11%), respectively. In general, the addition of Ti₂CT_x nanosheets made EP display better creep resistance and lower unrecovered strain than pure EP to a great extent at different stress levels.

Fig. 8 Creep and recovery of pure EP and its composites with different contents of Ti₂CT_x under different applied stresses of 5 MPa (a), 10 MPa (b), and 15 MPa (c) at 75 °C and (d) creep strain and unrecovered strain measured at 15 MPa and 75 °C.

The above-mentioned differences in the creep and recovery behaviour of composites with different contents of added Ti₂CT_x may be explained by several reasons. Firstly, the dispersion of a filler in the matrix has a major influence on the creep behavior of polymer composites. Ti₂CT_x particles with a layered structure were exfoliated and intercalated and distributed in the matrix evenly, establishing a close contact with EP molecular chains, and formed a larger interfacial area to improve the interfacial bonding and compatibility.^61,62 Secondly, Ti₂CT_x nanosheets interacted with EP molecular chains and occupied part of the free volume in the EP matrix, so that the molecular segments underwent reductions in activity and mobility. It was difficult for the molecular chains to move and furthermore for the modulus and dimensional stability of the composite to be improved, resulting in decreases in creep strain and unrecovered strain.⁶³ In addition, EP molecular chains were combined with Ti₂CT_x nanosheets to form a more stable three-dimensional crosslinked network framework during the epoxy curing process. As shown in Fig. 6, the fracture surfaces of Ti₂CT_x/EP composites are rougher than that of pure EP, which confirmed that rigid Ti₂CT_x nanosheets that were intertwined with the polymer matrix produced a larger interfacial area, prevented the motion of the molecular chains and dominated the processes of interlayer sliding and the tremendous deformation of the substrate under an external force, resulting in less creep deformation and a lower strain rate in the composite than in pure EP.^64–66 In essence, below the glass transition temperature pure EP and its composite exhibit mainly elastic deformation and a certain extent of viscous flow under a certain level of creep stress. After the addition of Ti₂CT_x nanosheets to the EP matrix, the internal friction due to the motion of EP molecular segments would increase and the composites would deform and display longer relaxation times.⁶⁷ Therefore, the composites exhibited less creep strain and permanent deformation after recovery behavior.

3.3.3 Analysis of mechanical properties. Surface hardness is an important parameter for measuring the application performance of materials. Fig. 9 shows a curve of Vickers microhardness values with an increase in Ti₂CT_x content. It can be seen that the microhardness of EP composites decreased first and then increased with an increase in Ti₂CT_x content. When the concentration of Ti₂CT_x reached 1.0 wt%, the value was reduced from 23.8 Hv to 18.2 Hv, which is about 23.5% less than that of pure EP. This trend is most probably related to the fluidity of molecular chains on the surface of the epoxy polymer. With the addition of inorganic nanoparticles, some defects were introduced into the composites and increased the free volume in the EP matrix, along with increasing the fluidity of EP molecular chains on the surface of the composite, namely, the surface hardness decreased. However, the microhardness of Ti₂CT_x/EP composites increased as the Ti₂CT_x content continued to increase. When a certain amount of Ti₂CT_x was added to the EP composite material, this played the role of supporting loads and enhancing its ability to resist deformation under external stress.^68,69 However, the results obtained from surface indentation tests have large standard deviation errors that should be taken into consideration. Moreover, an increase in the filler content may introduce more defects, which have an adverse effect on the mechanical properties of the composites.

Fig. 9 Vickers microhardness of pure EP and Ti₂CT_x/EP nanocomposites. The inset shows indentation images observed by a metallography microscope of pure EP and the composite with a Ti₂CT_x content of 2 wt%.

Its values of impact strength and flexural strength can reflect the toughness of a material indirectly. Fig. 10 shows the changing trends in the impact fracture strength and flexural strength of Ti₂CT_x/EP nanocomposites. As shown in Fig. 10(a), the impact strength of EP was significantly increased with the addition of Ti₂CT_x nanosheets. On an increase in the Ti₂CT_x content to 1.0 wt%, the impact strength reached a maximum value of 17.8 kJ m⁻², which represented an increase of 76% compared with that of pure EP. Similarly, in Fig. 10(b) the flexural strength reached 98 MPa when the mass content of Ti₂CT_x was 1.0 wt%, which, compared with Ti₃SiC₂/EP composites with a Ti₃SiC₂ content of 30 wt% (ref. 2) and Ti₃AlC₂/EP composites with a Ti₃AlC₂ content of 40 wt% (ref. 70) represented an increase of about 30% and 36%, respectively. The increase in impact fracture strength and flexural strength may have the following reasons. Firstly, the dispersion of Ti₂CT_x nanosheets in the polymer matrix has a major effect. According to the SEM analysis in Fig. 6, Ti₂CT_x nanosheets were dispersed into the EP matrix evenly and the composite exhibited good interfacial bonding. When the composite was deformed under the action of an external force, the Ti₂CT_x nanosheets would deform. Because the stiffness of Ti₂CT_x nanosheets is much greater than that of the organic matrix, when Ti₂CT_x nanosheets undergo deformation under longitudinal tensile stress and transverse compressive stress, this will result in hysteresis relative to the substrate in the entire process of deformation.^52,53 Furthermore, they absorbed more energy and increased the impact strength and enhanced the toughness of the composites. In addition, interfacial interaction is the basis of two-phase stress transfer. Therefore, the strength of interfacial bonding has a direct effect on the ability of composites to resist deformation. During the fracture process of Ti₂CT_x/EP nanocomposites, the external force gave rise to interfacial debonding between the inorganic nanosheets and the organic matrix. Crack propagation encountered the rigid Ti₂CT_x nanosheets, which deflected the developing cracks. As the cross-section morphology shows in Fig. 6, the roughness of the fracture surface increased with an increase in Ti₂CT_x content, and the composite had a larger total fracture surface area and dissipated more energy during the course of impact fracture compared with pure EP, exhibiting greater impact resistance and toughness.¹⁷ Besides, the Ti₂CT_x nanosheets promoted the generation of shear yielding under an external force in the composites. Interfacial debonding in combination with shear yielding consumed a large amount of energy during deformation, and then the composites displayed higher strength.^53,71 Nevertheless, there is a limitation on the filler content. The larger the content of fillers added to the polymer matrix, the more defects will be introduced into the composites, exactly as Fig. 10(a) and (b) shows that the impact strength and flexural strength exhibited reducing trends when the content of Ti₂CT_x exceeded 1.0 wt%. The current result also supports the fact that the fracture toughness of nanocomposites will decrease when the nanofiller concentration exceeds a certain weight fraction and give rise to cracks that are propagated unstably, as reported in the previous literature.^53,72,73 Based on the above-mentioned analysis, it should be essential to take the appropriate amount of nanofiller added to the polymer matrix into consideration to improve the material performance.

Fig. 10 Impact strength (a) and flexural strength (b) of composites with different mass concentrations of Ti₂CT_x.

3.3.4 Analysis of anti-friction properties. As components, the anti-friction properties of polymers are essential in engineering applications. Here, we investigated the relationship between the friction coefficient, filler content and sliding velocity by using a dry sliding friction pin-disk tester. Fig. 11 shows the friction coefficients of Ti₂CT_x/EP composites with different Ti₂CT_x contents, which were measured under different friction conditions. Obviously, it can be seen that the friction coefficient is closely related to the friction velocity and the concentration of Ti₂CT_x. The friction coefficient of the epoxy composite with a 2.0% mass fraction of Ti₂CT_x decreased by more than twice that of pure EP at a sliding velocity of 0.3 m s⁻¹. However, at a lower sliding velocity of about 0.15 m s⁻¹, the friction coefficient of the composite with a 2.0% mass concentration of Ti₂CT_x decreased by 40%, which demonstrates an obvious difference from the decrease at a high sliding speed. That is to say, it is more favorable to improve the anti-friction properties of the EP composite in high-speed sliding conditions.

Fig. 11 Friction coefficient of pure EP and its composites obtained by dry sliding wear and the relationship of the friction coefficient with different contents of Ti₂CT_x in nanocomposites and the sliding velocity or distance.

Fig. 12 shows FE-SEM micrographs of the worn surfaces of EP with different contents of Ti₂CT_x nanosheets. It can be seen that there was severe wear damage on the worn surface of the pure EP matrix (Fig. 12), and the dominant wear mechanisms were fatigue and fracture wear owing to the detachment of bulk materials and formation of cracks.^74,75 Compared with pure EP, the worn surfaces of EP nanocomposites with an increase in the content of Ti₂CT_x nanosheets displayed substantive changes. The phenomenon of fracture damage was gradually reduced and the worn surface became smoother with an increase in Ti₂CT_x content (Fig. 12(b)–(e)), which means that the nanofiller can transfer stress effectively and prevent the initiation of cracks and further extension of fractures. From the SEM analysis of the worn surface, it can be inferred that in addition to fracture and fatigue wear, adhesion wear also contributes. During the frictional process, Ti₂CT_x nanosheets distributed in the composites gave rise to surface friction with the steel pins directly and generated a large amount of heat, which softened the worn surface. Furthermore, intercalated Ti₂CT_x moved freely and became embedded into the softened worn surface at high loads and constituted a friction transfer film at the interface, providing a certain lubricating effect.^2,43,74

Fig. 12 FE-SEM micrographs of the worn surfaces of EP with different mass concentrations of Ti₂CT_x: (a) pure EP, (b) 0.5 wt%, (c) 1.0 wt%, (d) 1.5 wt% and (e) 2.0 wt%.

4. Conclusions

A novel 2-D transition metal carbide MXene (nanolayered Ti₂CT_x) was synthesised by etching a MAX phase in the liquid phase and subsequently structurally homogeneous Ti₂CT_x/EP nanocomposites were prepared by in situ intercalative polymerization. As shown by XRD and SEM, EP was intercalated into the Ti₂CT_x interlayer and the exfoliated Ti₂CT_x nanosheets were dispersed evenly in the EP matrix. The results of DMA and TGA revealed that the incorporation of Ti₂CT_x nanosheets into the EP matrix significantly increased the T_g and E_s values and improved the creep resistance and recovery properties, but reduced the thermal stability of EP. Compared with those of pure EP, the T_g and E_s values increased by 15% and 38%, respectively. In addition, an increase in the applied level of creep stress increased the level of creep deformation of the composites, which displayed much lower levels of creep strain with an increase in the Ti₂CT_x content. Furthermore, the formation of an interface and a crosslinked network structure had an advantageous effect on the mechanical properties. Although the surface microhardness of the composites decreased slightly, the impact strength and flexural strength increased by 76% and 66%, respectively, when the content of added Ti₂CT_x was 1.0 wt%. In particular, tribological studies revealed a lower friction coefficient at a high mass concentration of Ti₂CT_x and a higher sliding velocity than that of pure EP. The worn surface of the Ti₂CT_x/EP composite was smoother than that of pure EP, which suggested that severe wear damage was mitigated and the Ti₂CT_x nanosheets had a positive effect on solid lubrication.

Acknowledgements

The authors gratefully acknowledge the support of the National Natural Science Foundation of China (No. 51205111, 51472075 and 51275156), Plan for Scientific Innovation Talent of Henan Province (No. 134100510008), Program for Innovative Research Team of Henan Polytechnic University (T2013-4), and Opening Project of Henan Key Discipline Open Laboratory of Mining Engineering Materials (MEM12-11).

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