Embedding two-dimensional graphene array in ceramic matrix

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Science Advances  23 Sep 2020:
Vol. 6, no. 39, eabb1338
DOI: 10.1126/sciadv.abb1338


Dispersing two-dimensional (2D) graphene sheets in 3D material matrix becomes a promising route to access the exceptional mechanical and electrical properties of individual graphene sheets in bulk quantities for macroscopic applications. However, this is highly restricted by the uncontrolled distribution and orientation of the graphene sheets in 3D structures as well as the weak graphene-matrix bonding and poor load transfer. Here, we propose a previously unreported avenue to embed ordered 2D graphene array into ceramics matrix, where the catastrophic fracture failure mode of brittle ceramics was transformed into stable crack propagation behavior with 250 to 500% improvement in the mechanical toughness. An unprecedentedly low dry sliding friction coefficient of 0.06 in bulk ceramics was obtained mainly due to the inhibition of the microcrack propagation by the ordered 2D graphene array. These unique and low-cost 2D graphene array/ceramic composites may find applications in severe environments with superior structural and functional properties.


In the past decade, the discovery of graphene has stimulated worldwide interests on its exceptional physics and properties (1, 2). Although graphene has excellent mechanical properties and electrical conductivity, the search for the “killer application” for graphene is still in progress (3). One major promising direction is to make composites that use graphene as fillers (4) inside them; the graphene layers are spatially isolated and may possibly maintain their two-dimensional (2D) state, and the spectacular properties of graphene layers could thus be obtained in bulk quantities. Meanwhile, the composites are mechanically strengthened by the graphene fillers and also endowed with multifunction, such as thermal, electrical, and shielding properties, etc. (1, 5, 6). However, the field of graphene-reinforced composites has not yet exhibited substantial progress, mainly due to the following three reasons. (i) The price of graphene is still too high to be affordable in industry applications; (ii) controllable dispersion and orientation of the graphene layers in the composites is extremely difficult, as there is strong van der Waals force between the 2D graphene, leading to agglomerates in the matrix; and (iii) the lack of bonding between graphene filler and the matrix results in a low efficiency of load transfer, leading to poorer mechanical properties (5).

While there are many reports in composites with polymer and metal matrixes, the ceramic matrix composites containing graphene fillers started timidly a few years ago (7) because of the extreme challenge of processing. The insulating, stiff, and chemically inert ceramics with high melting point could be a unique matrix for graphene fillers to access their 2D state and maximize their intrinsic properties in various severe environments, if the graphene layers could be well dispersed and oriented (8). In this paper, we put forward a universal and delicate strategy to uniformly disperse and assemble 2D graphene array into ceramic matrix. Completely different from the conventional processing, we do not fabricate graphene first but start from an ordinary and cheap industrial raw material, expandable graphite. By treating it in a microwave oven, the interlayer spaces between graphene layers are expanded. Liquid ceramic precursor can be intercalated into the interlayer space under vacuum by assistance of organic coupling agent. In this way, ceramics and graphene layers can be alternatively mixed and stacked in a molecular level, finally resulting in composites with homogeneously dispersed and ordered aligned graphene layers. Strong matrix-filler bonding was also found as a result of the usage of the coupling agent, which leads to efficient load transfer and excellent mechanical properties.

To demonstrate the performance of this unexplored ceramic/graphene composites, we study their tribology properties. With the increasing demand of tribological applications, ceramics exhibit great potential in industry owing to the high elastic modulus, hardness, high compressive strength, low density, and chemical inertness compared with conventional metal parts (9). This includes ceramic bearings, cutting tools, extrusion dies, valves, sealing rings, and cylinder liners, etc. Especially, ceramics could be the best suitable material in those tribological applications with high load and precision requirement or corrosive environment, such as turbopump in space engine, total knee joint replacement, and hip joint replacement in biomedical engineering (9). However, the applications of ceramics are highly restricted by their intrinsic brittleness, which leads not only to a high friction coefficient due to microcrack formation under stress but also to catastrophic breakdown or fracture in operation. Here, through the incorporation of ordered graphene array, we reduce the friction coefficient of ceramics to an unprecedentedly low value, which brings new opportunities in tribology applications. We also achieve stable crack propagation in these ceramics, which can prevent the catastrophic failure in service. In addition, the ceramics show self-lubricating performance, which is particularly interesting in vacuum or aerospace, where organic lubricants cannot be applied.


Processing strategy of the 2D graphene array/ceramic composites

Incorporating graphene into ceramic matrix proves particularly difficult because it should be compatible with the existing ceramic processing, which needs high-temperature sintering process for densification. The conventional methods always start from commercial or home-made graphene powder or suspension, whose cost is still too high to afford in industry now. The graphene powder or suspension is first mixed with ceramic powder, colloidal suspensions, or sol-gel precursor and then consolidated by sintering. In this process, the graphene fillers always agglomerate together as a result of van der Waals force interaction between them, which become large-scale defects in the ceramic matrix and deteriorate the mechanical properties. In this paper, we put forward a facile, simple, and controllable approach to embed 2D graphene array in ceramic matrix (Fig. 1; see details in Materials and Methods and fig. S1 for further information). In contrast to those conventional methods that use graphene as the raw material, we start from a widely commercialized product, expandable graphite. First, we converted expandable graphite into expanded graphite (EG) by microwave heating; during this process, the interlayer spaces between graphene layers are expanded tens to hundreds of times along the direction perpendicular to the basal plane of graphite. A liquid ceramic precursor was intercalated into the interlayer space under vacuum with the assistance of an organic coupling agent (Silane, KH570). The product was further homogenized by ultrasonic treatment to form colloidal dispersion. Through these steps, the ceramic precursor and the graphene layers were mixed homogeneously. Afterward, the ceramic precursor was hydrolyzed into hydroxides and forms plate-like composite powders with the graphene layers as seen in fig. S1G. In the subsequent evaporation process, the evaporation of solvents could lead to the alignment of the sheets of hydroxide/graphene through the force of surface tension near the surface of the suspension. The ceramic hydroxide/graphene sheets were finally deposited onto the flat bottom surface of the container, which also behaves as a template and leads to further alignment of the hydroxide/graphene sheets. This prearrangement process was confirmed in fig. S1H, where an ordered and layered microstructure was observed.

Fig. 1 A facile route to synthesize the few layer graphene (FLG)/ceramic composite.

The ceramic/graphene precursor was then calcined to remove organics and then loaded into graphite die for spark plasma sintering (SPS), in which a certain degree of orientation disorder could occur. As the pressure was applied in the sintering process, the plate-like composites were gradually aligned with increasing temperature as clearly shown in fig. S2. At the beginning, the few-layer graphene (FLG) was not perpendicular to the pressure, and the shear force component along the layers caused sliding between the weak van der Waals interface inside the graphene layers. This could cause deflection of the graphene layers in combination with the flow and plastic deformation of the ceramic powder under the pressure at higher temperature. This deflection of graphene layers would stop until they were completely perpendicular to the direction of the pressure, when the shear force component along the layers becomes zero. Therefore, after SPS sintering, all the graphene sheets are all aligned along the direction perpendicular to the pressure.

Last, a ceramic matrix composite containing a graphene array was obtained. This method is easy to realize in industrial scale and is almost universal to all ceramic materials. We have fabricated different ceramic matrix composites including SiO2, Al2O3, ZrO2 (adding Y-stabilizer, aiming 3 mol % of Y2O3 in ZrO2 matrix), MgO, ZnO, TiO2, and Fe3O4, and all the samples exhibited a similar microstructure. This process is superior in economic cost. In contrast to most approaches that use expensive graphene as the raw material, we use expandable graphite instead, which has a typical price of $1 to $2/kg, nearly the same as that of normal bulk graphite. Therefore, the universality and low cost of this method could have high potential in future industry applications.

Microstructure characterization of the 2D graphene array/ceramic composites

The sintered composites were broken off, and the cross-sectional microstructure is shown in Fig. 2. The paralleled graphene sheets can be easily distinguished as the platelets protruding out from the fracture surface. Moreover, the lateral dimension of graphene is relatively large (~20 μm on average). The high aspect ratio makes the graphene layers efficient load carriers and makes it easier to arrest crack propagation, resulting in good mechanical properties. We also successfully fabricated MgO and functional ceramics and ZnO-, TiO2-, and Fe3O4-based composites using the same strategy, whose microstructures are shown in fig. S3 (A to D). Unlike SiO2, Al2O3, and ZrO2, these ceramic composites may show interesting functional properties, such as catalysis and quantum capacitance, which merit further investigation.

Fig. 2 Morphological and structural characterization of FLG/ceramic composites.

(A) Scanning electron microscopy (SEM) image of the fractured surface of 5 volume % FLG/SiO2. (B) HRTEM image of FLG in an SiO2 matrix. (C) X-ray photoelectron spectroscopy (XPS) spectrum for Si2p in the FLG/SiO2 composite. (D) SEM image of the fractured surface of 5 volume % FLG/Al2O3. (E) HRTEM image of FLG in an Al2O3 matrix. (F) XPS spectrum for Al2p in the FLG/Al2O3 composite. (G) SEM image of the fractured surface of 5 volume % FLG/ZrO2. (H) HRTEM image of FLG in a ZrO2 matrix with inset image showing the presence of dislocation near the interface. (I) XPS spectrum for C1s in the FLG/ZrO2 composite. a.u., arbitrary unit.

Because the three samples shown in Fig. 2 have the same volume fraction of graphene, the distances between the graphene layers are almost equal (fig. S3E). This further indicates controlled distribution of the fillers. The SiO2 matrix in this work is amorphous, while Al2O3 and ZrO2 have crystalline structures, as shown by the high-resolution transmission electron microscopy (HRTEM) images in Fig. 2 and the selected area electron diffraction (SAED) pattern in fig. S4. The scanning electron microscopy (SEM) observations of the samples reveal no obvious defects with pore-free and fine microstructures. At a higher magnification, graphene sheets with a considerably high aspect ratio can be distinguished (fig. S4D). In the crystalline matrix, graphene platelets are located both in grains and the interfacial boundary (fig. S4E). The HRTEM images show more details for the graphene layers (HRTEM images in Fig. 2). As illustrated in Fig. 2 (B, E, and H), a typical graphene platelet has six to eight layers of graphene, and the interlayer distance is approximately 0.5 nm, which is slightly larger than the spacing of graphite (0.334 nm) (10, 11). Because a single platelet with more than 10 layers of graphene has never been observed in the samples, we denote the composite in this work as FLG/ceramic composites. The graphene filler and the ceramic matrix are tightly combined with each other with a visible boundary. Especially, in the FLG/Al2O3 and FLG/ZrO2 composites, dislocation in the ceramic matrix near the matrix-filler boundary could be found (inset of Fig. 2H), which could be another evidence for the strong matrix-filler bonding. The chemical bond–coupled interface between the graphene and matrix was further confirmed by the x-ray photoelectron spectroscopy (XPS) of the three sintered FLG/ceramic composites, as shown in Fig. 2 (C, F, and I). In the FLG/Al2O3 composite, for instance, the Al─O─C bond at the interface between the graphene filler and the Al2O3 matrix can be detected (Fig. 2F). Therefore, in contrast to previous studies, the FLG/ceramic here shows strong matrix-filler bonding, which is beneficial for a high efficiency of load transfer.

FLG was further investigated by Raman spectroscopy analysis. As shown in fig. S5A, typical peaks were observed at ~1350 cm−1 (D band), ~1585 cm−1 (G band), and ~2700 cm−1 (2D band), which confirm the structure of FLG in the ceramic matrix (6, 12). Compared with that of the mixed powder, the intensity ratio of the D band and G band sharply increased after SPS densification. The elevated ID/IG ratio can represent the reduction degree of graphene (13). Usually, the ID/IG ratio is also used as a measure of carbon lattice disorder, which is expressed by the sp3/sp2 carbon bonding ratio (14). An increase in the ID/IG ratio suggests an increase in carbon disorder, which can also be seen from the HRTEM images in Fig. 2. The defective nature of FLG in the composites may be due to the chemical bonding with the ceramic matrix. The relatively lower G band in the composites may correspond to the deflection and wrinkling of graphene sheets (15, 16).

Ultralow friction coefficient of the 2D graphene array/ceramic composites

Ceramics have high melting point, chemical inertness, and compressive strength and are supposed to be competent in many severe friction applications, such as sealing and bearing parts at high temperature or in corrosive environments, cutting tools, and space airframes (17, 18). However, the friction coefficients of ceramics are rather disappointing, with typical values of 0.5 to 1.0. Ceramics are normally weak in tensile strength, and microcracks can be easily generated by the tensile force during friction. In addition, these cracks easily propagate because of the low toughness of ceramics and form loose wear particles when they intersect (17, 18).

To solve this severe issue, adding lubricating additive into ceramics has become a major route to reduce the friction coefficient, such as metals (19), oxides (2022), carbon nanotubes (CNTs) (2328), or graphene (26, 2939). These lubricating additives could form lubricating film on the surface of the ceramic and alleviate the microstress in friction, so the cracks and microfracture could be inhibited. However, until now, the friction coefficient can only be reduced by 10 to 40%, and most of them are still above 0.3. Here, we applied FLG arrays in a ceramic matrix and achieved notable improvements in the tribological properties, which were tested using a ball-on-disk configuration in rotary mode to the polished fracture surface (Fig. 3A) under different load and line speed. Figure 3B shows the friction coefficients of 5 volume % FLG/ceramic composites. The minimum friction coefficients of FLG/SiO2, FLG/Al2O3, and FLG/ZrO2 are extremely low, only 0.12, 0.06, and 0.06, respectively, which shows a huge reduction compared with the monolithic counterparts (fig. S6A). Figure 3C plots the friction coefficients for the different ceramics with the addition of graphene, CNTs, oxides, and metal fillers in addition with the monolithic ceramics. It could be found that most of the ceramics and ceramic-based composites show a friction coefficient above 0.3, and only a few can reach 0.2 (40). In contrast, the values of FLG/ceramic composites have exceeded those of the above composites and are almost insensitive to the tribological conditions. Especially, a new record for dry sliding friction coefficient of ceramics, 0.06, has been obtained in the FLG/ZrO2 composite.

Fig. 3 Tribological properties of FLG/ceramic composites.

(A) Schematic of the self-lubricating effect during the friction process and the SEM images of the wear tracks of the FLG-reinforced composite. (B) Friction coefficients of prepared FLG/ceramic (5 volume %) composites tested under different load and speed. (C) Comparison of the friction coefficient reduction in this work and that reported for ceramics and ceramic-based composites.

Mechanism for the ultralow friction coefficient

One of the mechanisms accounting for the ultralow friction coefficient is the formation of lubricating film due to the graphene filler. To verify its presence, SEM observation of the morphologies of the worn surfaces was performed. In the FLG/ceramic composites, FLG was pulled out of the composites and contacted with the wear debris on the worn surface to form well-consolidated graphene films (Fig. 3A). The spread of the films on the contact interface reduced the friction because of the low shearing strength of graphene. In contrast, in the monolithic sample, rupture symbols with particle debris existed on the worn surface (fig. S6B), which created a high stress and large strain area at the subsurface (41) and led to a large friction coefficient.

Nevertheless, the presence of lubricating graphene film alone cannot explain the ultralow friction coefficient of the current FLG/ceramic compounds compared with the other composites with varieties of lubricating additives in literature (Fig. 3C). The major reason could be related with the microfracture of the ceramics under friction. The FLG/graphene may have better resistance for the initiation and propagation of the microcracks as a result of improved mechanical properties.

Ceramics are much weaker in tensile strength than compressive strength. Cracks can be easily generated under the surface tensile force of friction. However, it is not always feasible to directly measure the tensile strength of ceramics. Instead, flexural strength is alternatively measured. Three-point bending tests were carried out to measure the flexural strength of the FLG/ceramic composites with an applied load perpendicular to the FLG filler, and the values are compared with those of the monolithic ceramics. Monolithic SiO2, Al2O3, and ZrO2 have flexural strength values of 65 ± 5, 424 ± 8, and 390 ± 8 MPa, respectively. Incorporation of 5 volume % FLG array increases the values to 99 ± 5, 560 ± 6, and 510 ± 5 MPa, respectively, and its increase approaches 30 to 50%, as shown in Fig. 4A.

Fig. 4 Mechanical properties and toughening mechanisms of FLG/ceramic composites.

(A) Flexural strength of three ceramics without and with FLG reinforcement (5 volume %). (B) Increasing crack extension resistance curves (evaluated by the steady-state fracture toughness, KJC) of three FLG/composites. (C) Crack deflection observed in in situ three-point bending test on the FLG/SiO2. (D) Microcrack deflection and graphene bridging in FLG/SiO2. (E) Wavy graphene in the ceramic matrix. (F) Stress distribution near the crack tip by progressive interface failure via finite element method simulation. (G) Comparison diagram of FLG/ceramic composites and other reported ceramics reinforced with carbon nanostructures.

The flexure strength of the FLG/ceramic composites closely depends on the microstructure and interfacial bonding between FLG and the ceramic matrix. Graphenes with a high Young’s modulus (42) strength and large specific area are well dispersed in the ceramic matrix, transferring the load from the matrix, which notably improves the composite flexural strength (4246). The homogeneous, parallel alignment of FLG works via a mechanism similar to that of “rebar” in “concrete.” Besides the increased bending strength, the composites have maintained the hardness of the monolithic ceramics (fig. S6C). Stiffness is also increased (fig. S6D) because of the strong matrix-filler bond and high efficiency of load transfer, which is beneficial for the reduction in friction coefficient (47).

While the increased bending strength (tensile strength) accounts for the resistance against the crack initiation under friction, the improved toughness is the major reason for preventing the propagation of these cracks. Single-edge V-notched beam (SEVNB) tests were performed to measure the mechanical toughness of the monolithic ceramics and the FLG/ceramic composites. Monolithic SiO2, Al2O3, and ZrO2 show KIC values of 0.73 ± 0.06, 3.41 ± 0.07, and 3.61 ± 0.07 MPa·m1/2, respectively. Incorporation of the ordered FLG array elevated the KIC values to 2.52 ± 0.06, 7.39 ± 0.05, and 7.44 ± 0.08 MPa·m1/2, respectively. Roberts (48) observed the minimum load required for abrasion-induced fracture for brittle materials based on indentation fracture mechanicsP=5447βπη2θ4(KICH)3KIC(1)where P is the minimum applied load (N) required to produce abrasion-induced surface fracture from a point contact, η is a constant, β is the constant relating hardness (2.16 for Vickers indentation), θ is the geometrical constant (≈0.2), KIC is the fracture toughness (MPa·m1/2), and H is the hardness (GPa) of the material. Above this load, there will be a transition from mild wear to severe wear, where the generation of crack and particle pullout can lead to an increase in friction coefficient. Taking Al2O3 for example, the mechanical properties (see table S2) of the samples with and without FLGs was incorporated into Eq. 1. It has been found that a minimum of ~2.2 N load is necessary for the abrasion-induced fracture of the investigated monolithic Al2O3 surface. We measured the friction coefficient under a quite broad range of loads, 5 to 30 N, which have exceeded the minimum load, which means monolithic Al2O3 enters the severe wear regime and leads to a high friction coefficient. In contrast, for the FLG/Al2O3 composites, the minimum applied load required for abrasion-induced fracture is as high as ~52.7 N, which is 20 times higher than that the value of monolithic Al2O3. Hence, the surface fracture and subsequent debris formation are notably suppressed. To prove this, we checked the surface morphology of the Al2O3 samples without and with FLGs (5 volume %) after friction test of the same conditions. As shown in fig. S6E, the friction surface of monolithic Al2O3 shows many microcracks and particle pullout. This substantially increases the surface roughness and leads to a large friction coefficient. As shown in fig. S6F, the surface of FLG/Al2O3 composites is much smoother than the monolithic Al2O3 and no microcrack can be seen, which accounts for the much lower friction coefficient compared with Al2O3. In addition, we notice that the grain size in the composites was refrained by the graphene array, which can also contribute to the low friction coefficient by inhibiting the fracture of grain boundary and grain pullout.

The graphene array modified the fracture behavior of the matrix. In the monolithic ceramics, linear crack extension leads to a catastrophic failure. In contrast, in the composites with FLG array, the crack is deflected by the interfaces and slowed down by various extrinsic toughening mechanisms, resulting in stable crack growth. This toughening leads to an increase in the fracture resistance as the crack propagates, known as the R-curve effect. To measure the R-curve, the indirect crack length is measured by a compliance method. The R-curves of all three FLG/composites are plotted in Fig. 4B. The curves exhibit an increasing behavior with a steady rise in KJC. The maximum toughness values of the different as-prepared FLG/composites were extremely high, approaching 4.21 ± 0.05, 12.43 ± 0.04, and 14.50 ± 0.06 MPa·m1/2. These values correspond to 500, 240, and 300% increase over the toughness of the monolithic ceramic, respectively. These values far exceed those of other carbon-reinforced ceramics (38, 42, 4960). In terms of energy, the critical strain energy release rates for the brittle monolithic ceramics SiO2, Al2O3, and ZrO2 are 8.9 ± 2, 86.6 ± 3, and 91.7 ± 4 J·m−2, respectively. The corresponding work of fracture is measured to be 76.2 ± 2, 225.9 ± 6, and 269.7 ± 5 J·m−2, which are 750, 160, and 190% improvement than the monolithic ceramics, demonstrating much higher resistance for propagation of microcrack under friction.

To explore the fracture process and clarify the toughening mechanism of our composites, an in situ three-point bending test was carried out using SEM to observe the crack evolution. The crack propagation of the graphene-reinforced composites exhibited a confluence of multiple toughening mechanisms (Fig. 4, C and D). In all three composites, the primary crack developed with a serpentine morphology (Fig. 4C and fig. S7, A and B) instead of a straight fracture (fig. S7C). The edge of the primary crack displayed an apparent zigzag path. Moreover, along with the primary crack, secondary microcracks occurred and propagated parallel to the plate (perpendicular to the propagation of primary crack; red arrows in Fig. 4D). All these mechanisms inherently elongate the crack length, absorbing more energy as the cracks propagate. In addition to the crack elongation and deflection, the Cook-Gordon toughening mechanism occurs in the composite (61). When a composite has soft layers (FLG) embedded within a hard matrix (ceramic), as a crack reaches the weak interface inside the FLG (fig. S7D), the stress on the crack can easily break the interface, forming a pronged crack ahead of the crack tip (62). In the composite samples, the FLG is perpendicular to the crack propagation direction; hence, the crack has to “break down” the ceramic matrix in a “step-by-step” manner. Thus, superparallel FLG arrays are crucial for composite sample toughening, as they maximize the Cook-Gordon effect. Two other toughening mechanisms are “graphene bridge” and “nanocrumpling.” Graphene bridges behind a crack tip inhibit opening of the crack (Fig. 4D), further preventing or delaying catastrophic fracture (63). Besides, the effect of loading rate on the composite’s toughness was considered. Three-point bending tests under different loading rates, 0.01 mm/min, 0.5 mm/min, and rapid shock test (about 0.5 to 0.8 m/s, to simulate the real service condition), were carried out to observe the crack propagation. Crack deflections were observed in all loading rates, as shown in fig. S7 (A, E, and F). The values of KIC of FLG/ceramics were also measured under different loading rates (0.05, 0.1, and 0.5 mm/min), as shown in table S1. It can be concluded that the loading rate has a very slight effect on the KIC values of our FLG/ceramic composites. In addition, a close observation of the fracture surface of the FLG/ceramic composite sample showed that the shape of the graphene filler is wrinkled (Fig. 4E), which can strengthen the mechanical interlocking with the ceramic matrix, preventing large-scale delamination (64, 65). Finite element method (FEM) simulation shows distinct microcrack deflection by progressive interface failure in the FLG array structure (Fig. 4F). In addition, FEM provides more insights into the toughening mechanism. In a monolithic ceramic, the maximum stress is supposed to be located at the crack tip in fracture. However, in these ceramic composites, the maximum stress is located at the graphene array, and the crack tip shows a much lower stress. This is because of the large stiffness and ordered orientation of the graphene fillers and their strong bonding with the ceramic matrix. The simulation result suggests that the strain field is redistributed by the graphene array in fracture, which provides a crack tip shielding effect and contributes to the high toughness.

At the sintering temperature, the ceramic matrix of SiO2 and ZrO2 could undergo phase transformation, which affects the microstructure and properties. The phase compositions of the sintered samples were analyzed by x-ray diffraction (XRD). For the SiO2 matrix composites, we can confirm that silica maintains an amorphous state, and no detectable crystallization was observed in fig. S8A. This result was consistent with the transmission electron microscopy (TEM) result (Fig. 2B) and the SAED pattern (fig. S4A). The amorphous structure was also observed in other reports of graphene/silica composites (56, 66). We also checked the XRD pattern of monolithic SiO2 after SPS sintering (fig. S8B), which suggests an amorphous structure. The silica was fabricated by hydrolysis of tetraethyl orthosilicate (TEOS) and heat treatment, which has an initial amorphous state. The powder was sintered at SPS for a very short time (4 min), and the amorphous structure was maintained (fig. S8B). Therefore, in the graphene/silica composites, the amorphous structure was maintained in the same way. In addition, it has been realized that the carbon nanostructure could further inhibit the crystallization of silica at high temperature due to the decreased viscosity of silica and lower mobility of silica network bounded to the surface of the carbon nanostructure (67). For the ZrO2 matrix composites, we confirm that ZrO2 has a tetragonal phase in fig. S8C. We have added 3 mol % Y2O3 to ZrO2 (3YSZ) to stabilize the tetragonal structure to room temperature, which has a critical grain size of 1 to 6 μm depending on the processing conditions (68, 69). We also did an SEM observation of the fracture surface and checked that the grain size of 3YSZ in the composites is around 200 nm (fig. S8D) because of the confinement effect of the parallel graphene array. This value is quite below that of the critical grain size, which accounts for its stable tetrahedral structure at room temperature. In contrast, for the monolithic 3YSZ, the grain size is around 6 μm (fig. S8F), which is close to the critical grain size. Therefore, the phase composition is mainly composed of the tetrahedral phase mixed with a few percent of monoclinic phase (fig. S8E). As is known, the stress-induced phase transformation in YSZ could enhance the mechanical toughness, and the effect depends on the difference between the real grain size and the critical grain size. In the graphene/3YSZ composite, the grain size is quite below the critical grain size, so the 3YSZ matrix in the composite has less tendency for phase transformation under stress and could have a lower toughness than the monolithic 3YSZ with a real grain size close to the critical grain size (70). However, the graphene/3YSZ composites have much higher toughness than the monolithic 3YSZ, demonstrating the extraordinary toughening effect of the ordered graphene array.

Bending strength and toughness are generally considered to be mutually exclusive (71). Because intrinsic toughening mechanisms are linked to plasticity and strength, a compromise is always reached in structural materials, and either of the properties may be sacrificed. In a previous work (72, 73), for instance, the introduction of graphene substantially increased the toughness but resulted in a notable decrease in the flexural strength. The mechanical properties of the monolithic ceramics and FLG-reinforced composites fabricated in this work are summarized in table S2, which shows remarkable improvement in both strength and toughness. We also compared the values with those ceramic-based composites reported in the literature (Fig. 4G). It shows that the FLG/ceramic composites exhibit a combination of enhanced mechanical strength and toughness.

As described above, this arises from the homogeneous dispersion and ordered orientation of the graphene fillers and their strong chemical and mechanical interlocking with ceramic matrix, which have been rarely established in the literature (fig. S9). Eventually, the FLG/ceramic composites exhibit ultralow friction coefficients, as the unique mechanical properties inhibit microfracture under friction.

Multifunction of the 2D graphene array/ceramic composites

Besides the ultralow friction coefficient and high mechanical toughness and strength, the graphene array/ceramic composites are supposed to have more potential superior properties, such as electromagnetic interference (EMI) property, catalysis, field emission properties, and quantum capacitance, etc., due to the unique microstructure of the ordered alignment of FLG in ceramic matrix. Here, we demonstrate that a record-high EMI shielding value in ceramics can be obtained. Although metals are the most common EMI material, ceramic shows advantages of high strength and hardness, light weight, and corrosion resistance (7477), which have incomparable superiority in applications of severe working conditions, such as high temperature, highly corrosive environments, or even with high mechanical loads. However, typical engineering ceramics like SiO2, Al2O3, and ZrO2 are insulating and almost transparent for electromagnetic radiation. Conductive materials with chemical inertness have been incorporated into the ceramics to optimize the EMI value, such as CNT, graphite, and graphene, but the resultant values are still too low as shown in Fig. 5B (7684). This is especially challenging, as the secondary fillers may deteriorate the mechanical properties due to poor dispersion or high filler loading.

Fig. 5 EMI shielding effectiveness (SE) of FLG/ceramic composites.

(A) EMI SE of FLG/SiO2, FLG/Al2O3, and FLG/ZrO2 composites in the X-band. (B) Comparison of EMI SE values from this work and those of other carbon/ceramic composites. (C) SER and SEA of FLG/SiO2, FLG/Al2O3, and FLG/ZrO2 samples at 8.2 GHz. (D) Schematic of the proposed EMI shielding mechanism, in which the electromagnetic wave undergoes multiple internal reflections between the aligned graphene layers and is absorbed into the ceramic matrix.

The FLG/ceramic composites show record-high EMI values compared with the composites fabricated by conventional methods (Fig. 5, A and B). The SET values reach 36.6, 40.2, and 43.5 dB for the SiO2, Al2O3, and ZrO2 matrix composites in the X-band. A huge improvement has been achieved compared with the blank ceramics SiO2 (0.002 dB), Al2O3 (0.1 dB), and ZrO2 (3.7 dB). The SET values of FLG/ceramics have far exceeded the value in commercial product and can be used in applications such as cell phone (~20 dB) (81). The shielding mechanism of the FLG/ceramic composites can be understood in Fig. 5 (C and D). Shielding due to absorption (SEA) is the dominant mechanism, rather than reflection (SER), which is different from those composites containing CNTs (85). As a result of the large surface area of the 2D FLG and their ordered alignment, prominent absorptions are expected because of the multiple internal reflections of the electromagnetic wave between the massive parallel interfaces inside the composites, which led to a higher EMI attenuation and an absorption-dominant shielding feature in the composites. Moreover, the aligned conductive FLGs and the confined ceramic layers form numerous microcapacitors inside the composites and can increase the permittivity and finally increase the electromagnetic absorption. It has been shown that the complex permittivity at 10 GHz has increased from 3.9–1.0 × 10−4i, 9.1–3.0 × 10−4i, and 22.7–1.6 × 10−4i to 190–2.33i, 240–2.95i, and 260–2.64i for SiO2, Al2O3, and ZrO2, respectively. Therefore, the ultrahigh EMI value with excellent mechanical properties makes the FLG/ceramic composites of particular interest in microwave absorption applications in severe working environments.


In conclusion, we put forward a previously unreported strategy to engineer a parallel array of 2D graphene into ceramic matrix through chemical intercalation of ceramic precursor into low-cost expandable graphite. The ordered 2D graphene array transformed the catastrophic fracture mode of brittle ceramics into stable propagation behavior, with 250 to 500% increase in mechanical toughness and 30 to 50% improvement in mechanical strength. An unprecedentedly low friction coefficient of 0.06 was thus obtained among ceramics, mainly due to the enhanced mechanical properties. Combined with the mechanical reliability and self-lubricating performance, these composites exhibit great potential in those tribological applications under severe environments, such as vacuum, high load, high wear, or with corrosive agents. In addition, the 2D graphene array/ceramic composites with unique structure exhibit record-low EMI property, as well as broad potential multifunctions, such as catalysis, field emission properties, and quantum capacitance. This paper demonstrates a previously unreported avenue to access the advanced properties of individual graphene layers in bulk quantities by embedding 2D graphene array in insulating, stiff, and chemically inert ceramics matrix, which is of particular interest in applications in severe environments.


Main raw materials

The main raw materials were expandable graphite (160 to 50 N, GRAFGUARD, USA), ceramic precursor (TEOS, Alfa Aesar, 99.5%; aluminum ethylate, Alfa Aesar, 99.0%; zirconium n-propoxide, Aladdin, 99.9%; yttrium nitrate Aladdin, 99.9%, aiming to 3 mol % yttria in zirconia matrix), coupling agent (KH570, Qingdao, China), and hydrolysis agent (aqueous solution of ammonia, Acros Organics, 70 wt %).


Fabrication of FLG/ceramic composites. We prepared the FLG/ceramic composites using an infiltration-intercalation process in which ceramic matrix is mixed and attached to EG at the molecular scale. The preparation process comprises the following steps, and the step number (1 to 7) corresponds to the number (red color) in Fig. 1.

Step 1: Microwave heat

First, we prepared EG by microwave treatment, during which a large amount of interlayer component would be released from the expandable graphite in the form of gas, namely, the flash vaporization effect (86). Then, the graphite chip expanded tens to hundreds of times along the direction perpendicular to the basal plane, leading to the formation of EG, as fig. S1 (A and B) shows. Compared with traditional heating in a high-temperature furnace, our method only needs a microwave oven operating for 15 to 20s, which is a very convenient operation and energy-saving method.

Step 2: Vacuum infiltration

Second, the EG was suspended in a vacuum for 30 min to eliminate adhering air between the EG flakes to assist infiltration. The combination facility is shown in fig. S1C. Under vacuum condition, the mixed solution of ceramic precursor and coupling agent completely immersed the EG, and the precursor was intercalated into the layer space of EG (fig. S1D). This ensured a molecular-level mixing of the raw materials.

Step 3: Ultrasonic homogenization

Subsequently, the suspension was ultrasonically homogenized (fig. S1E).

Step 4: Hydrolyzation

Then, the uniform colloidal dispersion systems underwent hydrolysis. In the step of hydrolysis, the ceramic precursor is transformed to hydroxide with the addition of hydrolysate (fig. S1F).

Step 5: Evaporation-induced prearrangement

Next, the colloid of hydrolysate was placed in a vacuum drying chamber, and FLGs combined with ceramic hydroxide coprecipitated at the bottom of the container as the solvent evaporates within 12 hours. During evaporation and precipitation, the prearrangement process took place, and the SEM of the dry precipitate is shown in fig. S1H.

Step 6: Heat treatment

Then, the precipitate was heated at 600°C for 3 hours under argon atmosphere to remove the residual organics and solvent by a tube furnace. After this step, the precursor had gradually changed into ceramic.

Step 7: Spark plasma sintering

To get the dense bulk composite, an SPS process was carried out. The equipment used for sintering was an SPS system (LABOX-225, Japan). All the relative density of the sample prepared here are above 98% after sintering at dwelling temperature (1350°C for FLG/SiO2, 1400°C for FLG/Al2O3 and FLG/ZrO2) with 100°C/min heating rate and a constant applied pressure of 40 MPa. In the case of sintering pressure, lamellar or plate-like materials can be preferentially aligned perpendicular to the direction of the applied pressure, forming oriented microstructures.

Last, the dense ceramic matrix composite containing parallel graphene layers was obtained (Fig. 1 and fig. S1I), and the ceramic particles are uniformly bonded with the graphene layers (Fig. 2).

Microstructural characterization

The SEM pictures were taken on Pt-coated samples by a Supra 55 microscope and a ZEISS New Vision 40. TEM was carried out on a Cs aberration corrected FEI Titan 80-300 S/TEM operated at 300 kV. TEM specimens were prepared from the bulk composite sample by focus ion beam milling using a Helios Nano Lab 600 instrument (2 to 30 keV Ga+ incident beam energy with currents of 16 to 21 nA). Laser scanning confocal microscopy (LSCM) photo was obtained on uncoated samples by a laser scanning confocal microscope (Olympus OLS41).

Geometric data analysis

The image analysis software (Image Pro Plus6) was used to collect the spacing between adjacent FLG and to get statistics from more than 80 data points.

X-ray photoelectron spectroscopy

High-resolution XPS spectra were measured with an ESCA Probe P from Omicron Nanotechnology (ESCALAB, Thermo Fisher Scientific, USA). Carbon conductive tape homogeneously covered with graphene was used for the measurements. An Al x-ray source with a monochromator was used for the excitation (fig. S5, B to D).

Raman spectrum

Raman spectrum of the polished surface perpendicular to the SPS-pressing direction was carried out by confocal micro-Raman spectroscopy. Raman maps of 30 × 30 pixels, recording one spectrum per pixel and using 1 s of acquisition time, were acquired on 30 × 30 μm2 scanned areas using a excitation laser wavelength of 530 nm.

Friction coefficient measurement

A universal tribo-tester (rotating end-face friction type, UMT-5, BRUKER, USA) was used to evaluate the friction behavior of the composites on the polished surface (perpendicular to the orientation of FLGs; roughness, <0.2 μm) at room temperature. The experiments were conducted at turning radius of 4 mm, rotating speed of 40 to 120 r/min (line speed of 8.4 to 25.1 mm/s), and constant load of 5 to 30 N, for the test duration of 60 min. The disks were Si3N4 balls of 3-mm diameter. The disks were polished, and average surface roughness was less than 0.1 μm.

Mechanical performance testing

The mechanical properties of the samples were measured by using a universal testing machine. Square-shaped samples with length of side of 20 mm (fig. S1I) were obtained after sintering. Beam-shaped specimens around 20 mm by 4 mm by 2 mm (toughness test) and 20 mm by 4 mm by 3 mm (strength test) were then cut from the sintered samples. The beams for the SEVNB (single-edge V-notched beam) test specimen were first notched with a diamond saw of 200-μm thickness, and then the bottom of each notch was sharpened by repeatedly passing a razor blade with diamond paste (1 μm). Using this method, the final notch radiuses were always below 20 μm (fig. S10A). Bending test was mirror polished and beveled to avoid any crack departure from the sides. Flexural strength was determined by a three-point bending test carried out on unnotched beams. At least eight specimens were tested for each composition.

For the R-curve measurements, samples were tested in four-point bending with a universal testing machine at a displacement rate of 0.01 mm/min. The samples were loaded until crack propagation was observed in the load/displacement curve. Afterward, the specimen was unloaded, and the crack was measured with LSCM (Olympus OLS41). Different measurements of crack propagation were taken with the precaution of loading the sample always in the same position. The deflections were measured by a linear variable differential transformer.

Fracture toughness, KIC, was calculated using following equations (87):KIC=PSBW32f(aW)(2)f(aW)=3(aW)12[1.99(aW)(1aW)(2.153.93aW+2.7(aW)2)]2(1+2aW)(1aW)32(3)in which P is the maximum load in the SEVNB test, S is the support span, B is the thickness of the specimen, W is the width of the specimen, and a is the notch depth. Fracture toughness, KJC, in this work was calculated from the elastic and plastic contribution, which relates to the J-integral calculation, similar to previously reported composites (8890)J=Jel+Jpl(4)

Jel is the elastic contribution, which is based on linear elastic fracture mechanicsJel=K2E(5)

The plastic contribution Jpl can be calculated with the following equationJpl=1.9AplB(Wa)(6)

Apl is the plastic area underneath the load-displacement curve. Thus, J values can be transformed into K values by the following equationKJ=(JE)12(7)where, E = E(1 − ν2), E is the Young’s modulus (obtained from fig. S10, B to G), and v is the Poisson’s ratio. As the variation of E influences KJC in a fairly limited way, here E′ can be replaced by E.

X-ray diffraction

XRD patterns were collected in step-scanning mode with Cu-Kα radiation in a range of detection of 10° < 2θ < 90°, with a scanning rate of 2°·min −1 in steps of 0.05°.

Finite element analysis

A 2D finite element (FE) model is developed using the commercial software ABAQUS v6.13. In the simulation, a 2D graphene array structure (10 × 4 mm2) with a single-edge notch (2 mm by 0.02 mm) is adopted (see fig. S10H). This structure in the FE model contains the inerratic arrangement of graphene sheets bonded with the ceramic matrix, which is modeled as a cohesive zone with a bilinear traction separation. The interface between two parallel ceramic brick is also modeled as a cohesive zone with a bilinear traction separation. The graphenes with isotropic bulk modulus Ep = 1 TPa, Poison ratio vp = 0.16, and failure strength σmp = 100 MPa undergo elastic deformation before brittle failure. In addition, the ceramic matrix with isotropic bulk modulus Ep = 72 GPa, Poison ratio vp = 0.23, and failure strength σmp = 10 MPa undergo elastic deformation before brittle failure. The initial response of the cohesive element is assumed to be linear until a damage initiation criterion is met. The energy-based damage evolution criterion with mode-independent fracture is adopted. When the accumulated energy release rate G is larger than the critical energy release rate Gc, the interface is fully fractured. The critical energy release rate Gc are 1.2 and 0.89 N m−1 for graphene and ceramic, respectively. We simulated the crack propagation with the three-point bending. The points away the ends of the bottom 2 mm are fixed, and loading is applied at the center of the top. We have chosen the parameters in the ABAQUS model such that crack propagates in our numerical three-point bending test shown in fig. S10I and can recover to the similar experimental three-point bending.

Electromagnetic shielding property measurement

For the EMI shielding effectiveness (SE) dielectric constant characterization in the X-band frequency (8.2 to 12.4 GHz), 22.86 mm by 10.16 mm by 2.00 mm specimens (fig. S1I) were polished and the S-parameters (S11 and S21) of each sample were determined in the X-band through the waveguide method with a vector network analyzer (Agilent N5230A). For accuracy of measurement, the device is carefully calibrated with the through-reflect-line (TRL) approach. A MATLAB code based on the S-parameters was developed to extract shielding by reflection, shielding by attenuation, and the total SE. The major EMI shielding mechanisms include reflection, absorption, and multiple reflections (16, 91, 92). Reflection arises from the interaction between mobile charge carriers and the electromagnetic fields; absorption loss results from interactions between electric and/or magnetic dipoles and the electromagnetic fields. Therefore, the total EMI SE (SET) of a shield can be expressed asSET=SER+SEA(8)where SER, SEA, and SEM denote the shielding effectiveness due to reflection, absorption loss, and multiple reflections, respectively. Experimentally, SER and (SEA + SEM) can be calculated as follows (85)SER=10lg(1S112)(9)SEA=10lg(S2121S112)(10)where S11 and S21 were normalized S-parameters that were obtained from the vector network analyzer.


Supplementary material for this article is available at

This is an open-access article distributed under the terms of the Creative Commons Attribution-NonCommercial license, which permits use, distribution, and reproduction in any medium, so long as the resultant use is not for commercial advantage and provided the original work is properly cited.


Acknowledgments: Funding: This work was mainly funded by the National Natural Science Foundation of China (grant nos. 51590893 and 51788104) and the National Key R&D Program of China (grant no. 2017YFA0700705). L.Y. acknowledges financial support from the National Natural Science Foundation of China (grant no. 11890684). Author contributions: C.W. designed the experiments. C.S., Y.H., W.W., P.Z., and Y.X. performed the materials fabrication and properties measurement. Q.S. and L.Y. did the FEM simulation. C.S., W.P., and C.W. analyzed the data and wrote the manuscript. Competing interests: The authors declare that they have no competing interests. Data and materials availability: All data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials. Additional data related to this paper may be requested from the authors.

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