Hot-pressed graphene nanoplatelets or/and zirconia reinforced hybrid alumina nanocomposites with improved toughness and mechanical characteristics

1. Introduction

Fabrication of hybrid ceramic nanocomposites is a contemporary strategy in ceramic composites processing technology that promises better materials than those available today. The hybrid nanocomposite system can be produced by reinforcing the ceramic matrix with two or more nanostructured additives that possess distinctly different morphologies, compositions, and/or intrinsic properties. A superior mechanical response as well as unique functional properties are anticipated from the hybrid ceramic nanocomposites designed for high-performance, yet cost-effective commercial applications [1]. Furthermore, the new hybrid ceramic nanocomposites are expected to overcome the inherent brittleness issue associated with almost all the monolithic ceramics.

Alumina (Al₂O₃) is a promising ceramic for manufacturing of cutting tool inserts, electrical insulators, biomedical implants and as lining material for chemical containers. Although alumina stands out among the technical ceramics due to higher elastic modulus and hardness values, relatively low cost and abundance, its low fracture toughness and inferior ballistic properties limit its use in advanced armor, aerospace and automotive applications [2]. In case of layered ballistic armor system, the ceramic material faces the initial impact of the projectile to perform imperative functions of projectile speed reduction, blunting of the sharp geometry and conversion of the projectile into harmless tiny fragments [3]. For this purpose, the ceramic armor requires combination of high hardness and superior fracture toughness, the traits that monolithic alumina does not offer. However, microstructural tuning of alumina with nanostructured reinforcing phases is a potential approach towards production of high performance composite materials suitable for armor system.

Over the past two decades, dual-phase alumina ceramic nanocomposites have been made by incorporating carbon, ceramic, and/or metal based reinforcing nanomaterials with a single objective of toughness enhancement [4]. In this context, zirconia (ZrO₂) nanoceramic was particularly added to the alumina ceramic for preparing fine-grained dual-phase nanocomposites having improved corrosion resistance, electrical conductivity, biocompatibility, low coefficient of friction and higher wear resistance as well as superior strength for biomedical and other engineering applications [5,6]. It was found that the alumina nanocomposite containing 2-20 wt% zirconia nanoparticles demonstrated 25%-45% higher fracture toughness (K_IC) values as compared to monolithic alumina [7,8]. Such variations in the degree of improvement in mechanical properties was attributed to the challenges of zirconia nanoparticles distribution homogeneity in the alumina matrix, large range of grain size, surface and bulk porosity as well as poor process control during sintering [[9], [10], [11]]. In zirconia/alumina nanocomposite system, alumina contributes to the strength and hardness whereas zirconia ceramic increases K_IC by generating micron level cracks in the alumina matrix through a phase control transformation toughening mechanism [12].

Among carbon based nanomaterials, graphene is considered potential nanofiller for alumina ceramic owing to its two dimensional (2D) carbon sheet geometry arranged in a planar honeycomb lattice configuration that exhibits remarkable properties such as greater thermal conductivity and Young’s modulus values [13]. Alumina ceramic matrix reinforcement with up to 0.5 wt% graphene nanoplatelets was reported to cause an increase in the K_IC value by 73%, whereas greater graphene additions (>0.5 wt%) could barely lead to any further increment in the K_IC values [[14], [15], [16], [17]]. Although dual-phase alumina matrix composites showed improved mechanical performance, they could not fulfill the criteria required for use in the armor applications, thus leading to the shift in focus towards development of complex hybrid nanocomposite systems.

The hybrid nanocomposite design with essentially three reinforcing phases has evolved as one of the current research methodologies for producing tough, strong and hard ceramic nanocomposites. The co-additions of SiC and ZrO₂, or SiC and TiC nanoparticles into alumina was reported to yield hybrid nanocomposites with 95 and 30% higher K_IC values, respectively, as compared to monolithic alumina [18,19]. Other efforts employed spark plasma sintering and induction heat sintering techniques to produce alumina hybrid nanocomposites containing both fibrous MWCNTs and SiC nanoparticles with resulting increments in the K_IC values by 70% and 110%, respectively [20,21]. The coexistence of both SiC nanoparticles and graphene nanoplatelets as reinforcements in hybrid alumina ceramic nanocomposites, consolidated by means of induction heat sintering and spark plasma sintering techniques, resulted in higher K_IC (by 160% and 50%) and hardness values (by 27% and 36%) with respect to pure alumina, respectively [22,23]. Furthermore, spark plasma sintering of hybrid alumina nanocomposites, incorporating both CNTs and graphene nanoplatelets, exhibited increments in the K_IC and flexural strength values by 62% and 18%, respectively [24]. Hybrid reinforcing designed was also practiced for improving the performance of other ceramics in terms of K_IC values, elastic modulus, wear resistance, microhardness and creep properties [25,26]. Other instances of hybrid nanocomposites involved silicon nitride (Si₃N₄) and tantalum carbide (TaC) as ceramic matrices with CNT and TiN co-addition and subsequent respective improvements in the K_IC values by 45% and 200% [25,26].

Despite the advantages of hybrid nanocomposites, the studies on zirconia and carbon-based nanostructures (CNTs and Graphene) co-addition into the alumina matrix are scarce [27,28]. Besides zirconia nanoparticles, addition of MWCNT or graphene nanoplatelets caused a respective 30% and 45% enhancement in the K_IC values in the resulting hybrid alumina nanocomposites [27,28]. Keeping in view the potential applications of the hybrid alumina/zirconia/graphene nanocomposites for advanced aerospace and armor technologies, extensive efforts are imperative to exploit the outstanding properties of the GNP and zirconia nanoparticles. However, the contemporary challenges of uniform dispersion of the nanoscale reinforcing phases and strong interfacial bonding in the hybrid nanocomposites require immediate attention to produce high quality hybrid nanocomposites for the advanced armor as well as aforementioned applications.

In this work, hybrid alumina nanocomposites containing both zirconia nanoparticles and graphene nanoplatelets (GNP) were fabricated using a hot pressing machine. Monolithic alumina as well as alumina nanocomposites reinforced with zirconia nanoparticles (4 wt%) or/and GNP (0.5 wt%) were prepared following the same processing steps for comparison purposes. The consolidated nanocomposites were thoroughly scrutinized for densification behavior, dispersion and grain size, and mechanical properties using diverse analytical tools. The ballistic performance was gauged keeping in view the potential application of hybrid nanocomposites in the armor technology. Moreover, both qualitative and quantitative approaches were employed to assess the reinforcing capabilities of the GNP and zirconia nanoparticles in the hybrid alumina nanocomposite. The work also presents and discusses the reinforcing ability and the possible underlying toughening mechanism(s) activated by co-presence of the GNP and zirconia nanostructures in the hybrid nanocomposite.

2. Experimental techniques

2.1. Hybrid nanocomposites: powders mixing and consolidation

The GNP nanomaterial was synthesized by performing both oxidation and thermal treatments over graphite flakes (Asbury Graphite Mills Inc., NJ, USA). The processing route involved initial transformation of the graphite flakes (1.5 g) into graphite oxide (GO) using the protocol described by Hummer’s method [29] followed by distilled water rinse for cleaning purpose. An induction heating furnace (HF Active Sinter System, ELTEK, South Korea) at King Saud University, Saudi Arabia was operated up to 1550 °C at 1000 °C min^-1 heating rate under high vacuum environment of 45 mtorr for the GO thermal exfoliation process. The exfoliated GNP nanomaterial was thoroughly rinsed using distilled water to remove undesirable attached byproduct residues. Alumina (φ < 50 nm, melting point 2040 °C, specific surface area >40 m² g^-1) and zirconia (φ < 80 nm, melting point 2700 °C, specific surface area >25 m² g^-1) ceramic nanopowders with particulate morphology were purchased from Sigma Aldrich, UK. During hybrid nanocomposite mixing process, the alumina and zirconia nanoparticles were first homogenously mixed together employing an environmental friendly ultrasonic probe (Sonic Vibracell, VCX-750, Sonics, USA), to support colloidal chemistry process. The GNP nanomaterial was separately suspended into distilled water practicing a combined ultrasonication and wet chemistry process, according to the strategy mentioned elsewhere [30]. In a later step, the alumina/zirconia aqueous suspension was slowly added to the GNP aqueous suspension and the resulting mixture was further mixed through sonic vibrations for 60 min. The water from alumina/zirconia/GNP slurry was evaporated at 120 °C and the dried hybrid nanocomposite powder was obtained for onward hot-pressing. Following the aforesaid mixing process, monolithic alumina and nanocomposites reinforced with alumina and/or zirconia were consolidated with final compositions given in Table 1. The dried loose powder mixture was hot-pressed (S8538 UEO, FCT system, Germany) at the University of Exeter, UK. During hot-pressing, a 10 g dried powder mixture was filled inside a square hole inside a graphite die and then pressed at 50 MPa using graphite plungers. The hot-pressing was performed at 1600 °C under constant Ar atmosphere to obtain square-shaped solid samples with dimensions of 15 mm × 15 mm × 3 mm for all the samples.

2.2. Material characterization and fracture toughness investigation

The apparent densities for all the hot-pressed solid samples were measured by following the ASTM standard (C20) employing a machine (Sartorius Lab Instruments GmbH, Germany) working on Archimedes’ principle. For relative density calculation, the theoretical densities of alumina, zirconia and GNP were considered to be 3.97, 5.68 and 2.0 g cm^-3, respectively [22,31]. A pulse-echo method was adopted for determining the sonic velocity in all hot-pressed samples and for this purpose, a contact transducers (ϕ 6 mm) manufactured by TMP-3, Sonatest, UK was operated at fixed centre frequency of 20 MHz. The phase composition of the hot-pressed samples was appraised by means of an X-ray diffraction machine (D-8 Discover, Bruker, Germany) and CuKα monochromatic radiation was used. The microstructural features of all the hot-pressed samples were explored by operating a field-emission scanning electron microscope (JSM-7600F, JEOL, Japan). The grain size was estimated from the SEM micrographs according to the linear intercept method described in the ASTM E 112 standard. A high-resolution field-emission transmission electron microscope (2100F, JEOL, Japan) was used to characterize the intrinsic structure of the nanocomposites components and to investigate the GNS/alumina interfacial structure. The samples for TEM analysis were first dispersed in acetone and then transferred onto a perforated carbon-coated copper TEM grid. An ion-slicer machine (EM 09100, JEOL, Japan) working on the ion-milling technique was used to prepare ultrathin hybrid nanocomposite sample cross-section for interfacial characterization.

The microhardness values of all the hot-pressed samples were determined on a Buehler-Micromet 5114 hardness testing machine from Akashi Corporation, Japan. The microhardness was recorded on polished hot-pressed samples at a fixed applied load of 9.8 N for 15 s holding time. The nanoindentation experiment was conducted on the well-polished solid sample surface according to ASTM standard C1327-99. A typical pyramid-shaped Berkovich diamond tip installed in Hysitron Tribo950 machine at Royal Melbourne Institute of Technology, Australia was used for indenting the sample surface at specific parameters of load (1000 μN) and dwell time (10 s). The indents were carefully generated by maintaining an equal space of 20 μm in order to minimize the residual stress. From load-displacement curve, we carefully extracted the essential data regarding the displacement (h_max), applied force (F_max) and the elastic contact stiffness $[C=\frac{dF}{dh}]$ and computed the reduced modulus $[E_{r}=\frac{\sqrt{π}}{2C\sqrt {AP}}]$ values. The aforementioned parameters were further used for estimating the modulus of elasticity (E) through Oliver and Pharr’s model shown in Eq. (1) [32]. In Eq. (1), the V_s and V_i represent the respective values of Poisson’s ratio for the sample (0.23) and the Berkovich indenter (0.07); E_r and E_i are the reduced modulus values for the sample and the diamond indenter (1140 GPa), respectively [33]:

$E=\frac{(1-V^{2}_{s})}{\frac{1}{E_{r}}-\frac{(1-V_{i}^{2})}{E_{i}}}$ (1)

In this study, we practiced an established indent-crack method (ICM method) for estimating the K_IC values for all the hot-pressed samples. ICM is an efficient, yet experimentally simple method for gauging the K_IC values of the ceramic-based composites. This method is based on the length of the cracks that originate from an indent corner upon microindentation of the polished hot-pressed solid sample during Vickers microhardness test. For each sample, around 10 indents were made and the crack length values were carefully measured using SEM. Finally, the K_IC values were calculated by a mathematical relationship suggested in Chantikul model, as presented in Eq. (2) [34], where E, H, P and c represent the elastic modulus, Vickers hardness, the applied load, and the crack length, respectively:

$K_{IC}=0.016[\frac{E}{H}^{\frac{1}{2}}][\frac{P}{C^{\frac{3}{2}}}]$(2)

3. Results and discussion

3.1. Characterization and dispersions of the reinforcing inclusions

The graphene nanoplatelets (GNP) synthesized through thermal expansion route were characterized using SEM and TEM, as shown in Fig. 1. The SEM micrograph of the GNP (Fig. 1(a)) revealed two-dimensional (2D) sheet morphology with lateral surface length estimated to be in the range of 2-3 μm, and estimated thickness of ~5 nm (an upright GNP cross-section edge shown as Fig. 1(a) inset). The TEM analysis of the GNP also indicated nanosheet-like 2D configuration along with crumpling as well as overlapping features, as marked in Fig. 1(b). During rapid heating stage, the GNP undergo thermal expansion with associated stern chemical (intercalates eradication) and physical (graphene layers reshuffling) circumstances, in accordance with an earlier report [35]. The aforementioned shred of evidence from extensive electron microscopy sessions corroborates the strong GNP stability in preserving its extrinsic morphology as well as intrinsic structure during severe processing conditions.

The SEM analysis of the hybrid nanocomposite powder after colloidal mixing indicated presence of uniformly dispersed alumina and zirconia nanoparticles over the GNP lateral surface (white circle in Fig. 1(c)). The 2D geometry and large lateral surface area of the GNP has been reported to be beneficial towards attaining uniform dispersion of the ceramic nanoparticles, which is in contrast with the CNT behavior due to their strong tendency for entanglement owing to 1D fibrous geometry and high aspect ratio [36]. This elucidates the synergistic effect of GNP nanoscale morphology and highly energy sonic waves in dispersion and attachment of the alumina and zirconia nanoparticles onto the GNP surface via a colloidal suspension approach. The strong adhesion of the ceramics nanoparticles with the GNP (Fig. 1(d)) may be attributed to the phenomenon of acoustic cavitations that are produced in aqueous suspension due to powerful sonic waves. It is a known fact that a sonic wave can generate extremely high temperatures (~5000 °C) and pressures (~50 MPa) at the ceramic/graphene interface for a very short time period (microseconds) [37]. Hence, strong attachment of the alumina and zirconia nanoparticles to the GNP surface is certainly beneficial as it both sustains the dispersions of these nanoparticles during entire mixing process as well as prohibits the GNP folding during hot-pressing stage. The SEM examination of the hybrid nanocomposite fractured surface also confirmed effectiveness of the colloidal processing route in terms homogeneous distribution of the GNPs and zirconia, as shown in Fig. 2(a). Furthermore, the presence of all the reinforcing constituents in the hybrid nanocomposite was verified through presence of peaks characteristic of Al, Zr, O and C elements, during EDS analysis (Fig. 2(b)).

3.2. Microstructural development and phase analysis

For hot pressed monolithic alumins, A₄Z and A_0.5GNP nanocomposites and A₄Z_0.5G hybrid nanocomposite, the relative density values were calculated to be 99.7%, 99.4%, 99.2% and 98.5%, respectively, as listed in Table 1. The densification behavior of ceramic nanocomposites reinforced with nanostructured phases relies on annihilation/minimization of porosity and movement of grain boundary area via bulk diffusion phenomenon during high temperature consolidation [38]. However, the nanoscale reinforcement adversely affect the aforementioned densification phenomenon through segregation at ceramic matrix grain junctions, thus obtaining dense microstructure in the consolidated sample is challenging [39]. In our hybrid nanocomposite samples, the high densification is presumably due to uniform distribution of the nanostructured reinforcements in the matrix as well as simultaneous application of high pressure and temperature during hot pressing, both of which contributed to porosity reduction and mass transport phenomenon in the hybrid nanocomposite. In addition to that, the GNP also supported the consolidation process by: (i) acting as a solid state lubricant to the growing matrix grains at optimum adjusting position for minimizing the defects and porosity in the microstructure and (ii) enhancing heat flow in the powder compact owing to its excellent intrinsic thermal conductivity [40]. Table 1 shows that co-addition of zirconia nanoparticles and GNP led hybrid nanocomposite to have grain refinement by 68% as compared to monolithic alumina. On the other hand, nanocomposites incorporating zirconia nanoparticles or GNP demonstrated grain size reduction by only 20 and 35%, respectively. The microstructural refinement in the nanocomposites reinforced with either nanofiller is attributed to the uniform distribution of the nanofiller and their segregation at grain junctions during consolidation. Both mechanisms restrict grain growth through Zener grain pinning phenomenon, as observed in the hot-pressed sample (Table 1). Furthermore, immediately after fracture, the qualitative analysis of all the samples was performed using SEM for identification of the failure-mode as well as grain size determination, as showcased in Fig. 3. While large grain size with inter-granular failure mode was observed in pristine alumina (Fig. 3(a)), addition of reinforcing fillers resulted in grain refinement along with predominantly inter-granular failure mode, as displayed in Fig. 3(b)-(d).

Fig. 4 shows the XRD patterns of all the hot-pressed samples with diffraction peaks corresponding to the α-alumina phase (JCPDS No. 01-078-2426). The XRD pattern of the A_0.5G sample (Fig. 4(a)) indicated presence of an additional peak at 26° which was indexed to be the basal plane in crystalline graphite (JCPDS card no. 01-075-1621). The phase analysis of the A₄Z nanocomposite manifested addition peaks (Fig. 4(b)) which were assigned to the zirconia phase (JCPDS file no. 00-065-0728). The XRD profiles of the hybrid nanocomposite sample revealed diffraction peaks representative of the alumina matrix, zirconia nanoparticles and GNP fillers, as shown in Fig. 4(b). The XRD analysis also confirmed preservation of the reinforcing phase(s) composition/morphology in the nanocomposites despite harsh consolidation conditions.

3.3. Mechanical and ballistic properties

The K_IC value for each hot-pressed sample was determined using the ICM method, as tabulated in Table 2. The K_IC value for the hybrid nanocomposite (A₄Z_0.5G) was determined to be 7.9 MPa m^1/2, which is greater than those of the pure alumina (3.1 MPa m^1/2), A₄Z (4.6 MPa m^1/2) and A_0.5G (5.9 MPa m^1/2) samples by 155%, 72% and 38%, respectively. The evaluation of the K_IC value using ICM method is rendered debatable since this technique yields localized K_IC value. In contrast, the single-edged notched beam (SENB) method is considered a standard approach since it offers the bulk or absolute K_IC value [41]. In this context, a recent study about GNP/alumina ceramic nanocomposites reported insignificant difference between the K_IC values evaluated by ICM and SENB methods [17]. Nevertheless, ICM is an experimental, yet relatively simple approach to evaluate any improvement in the K_IC value by comparing it with the benchmark specimen. In theory, toughness is related to the higher critical energy dissipation rate, which is defined as the energy required for crack propagation [41]. Therefore, the critical energy dissipation rate (G_IC) for each sample was computed by a mathematical expression given in Eq. (3) [42] and the values obtained are given in Table 2.

$G_{IC}=K^{2}_{IC}(\frac{1-v^{2}}{E})$ (3)

where parameters K_IC, E and v represent the fracture toughness, the elastic modulus, and the Poisson's ratio (0.21), respectively [42]. For the hybrid nanocomposite sample (A₄Z_0.5G), the G_IC value was computed to be 265 J m^-2 which is greater than the A₄Z, A_0.5G and monolithic alumina samples by 38%, 85% and 191%, respectively. Such significantly higher G_IC value underlines the greater energy requirement for crack opening, thus interpreting superior resistance to crack propagation and forecasting a greater K_IC value. Additionally, the higher G_IC value leads to the anticipation that GNP and zirconia would activate distinct toughening mechanisms in the hybrid nanocomposite that will be discussed in Section 3.4.

From mircohardness data, the hybrid nanocomposite (A₄Z_0.5G) was found to exhibit ~17% and 10% increase in hardness as compared to monolithic alumina and A_0.5G samples, respectively. The higher microhardness values in case of A₄Z_0.5G and A_0.5G nanocomposites may be associated with GNP dual role towards microstructural refining and strengthening of the matrix grain boundaries. The greater lateral GNP surface area facilitates zirconia nanoparticles in attaining uniform distribution in the parent matrix with increased efficiency of alumina grain pinning by zirconia nanoparticles. In polycrystalline materials, grain refinement in the microstructure leads to an associated increase in the hardness as described by the Hall-Petch relationship that correlates hardness and grain size [43]. The fine-grained microstructure effectively obstructs the dislocations movement across the matrix grains, since grain boundary areas obstruct the onset of plasticity. Furthermore, the GNP induce unique mechanism for strengthening the matrix microstructure, a phenomenon that improves the load bearing capability of the GNP with subsequently greater hardness in the hybrid nanocomposites. A drop in the microhardness value by 12%, as in case of A₄Z nanocomposite, may be attributed to porosity and grain boundary cracking phenomena induced by phase transformation in the zirconia nanoparticles [44]. The A₄Z_0.5G hybrid nanocomposite, A₄Z and A0.5 G nanocomposites showed 13%, 11% and 8% lower elastic modulus values than that of monolithic alumina. The gradual drop in the elastic moduli of alumina upon incorporation of reinforcing phases may be attributed to microstructural defects and individual elastic modulus values of the reinforcing constituent phases [45]. The relatively low elastic modulus and microcracking phenomenon due to zirconia phase transformation from tetragonal to monolithic crystal structure are believed to have caused lowering of the elastic modulus in the A₄Z nanocomposite [45]. The large disparity in the GNP elastic moduli (in-plane and out-of-plan) also seems to be another reason behind the inferior modulus values in case of A_0.5G sample [46].

3.4. Reinforcing ability and toughening mechanism of the hybrid nanocomposite inclusions

Low and high magnification SEM views of the indent after microhardness test and fracture surfaces of the A₄Z and A₄Z_0.5G nanocomposites are shown in Fig. 5. The indentation appears to have cracks originating from its edges with certain length and pattern (Fig. 5(a)) in case of the hybrid nanocomposite (A₄Z_0.5G). The crack crevice was carefully explored at high magnification to develop an insight into the interaction between reinforcing nanostructures and the alumina matrix as well as for understanding the active toughening mechanisms in the hybrid nanocomposite. The in depth crack area analysis demonstrated bridging of cracked alumina surface by a GNP segment (Fig. 5(b)). A GNP segment partially detached from the matrix was also identified in the crack crevice (black arrows in Fig. 5(b)) with its other end firmly affixed with the matrix grain. The SEM analysis of the A₄Z nanocomposite microstructure exhibited well-dispersed zirconia nanoparticles at the matrix grains junction as manifested by black circles in Fig. 5(c) and boundary cracking (blue arrows in Fig. 5(c)). The presence of such microstructural features validates activation of grain boundary cracking toughness mechanism in the A₄Z nanocomposite. The expansion of zirconia nanoparticles is responsible for cracking at the zirconia/alumina interfacial region following transformation from the tetragonal to monoclinic phase [47]. According to the failure theory, the grain boundary cracking phenomenon in zirconia-reinforced alumina nanocomposites narrows the critical flaw to an extremely diminutive size at the notch tip area and this phenomenon leads to a substantial rise in the energy dissipation zone resulting in a greater K_IC value [45]. In contrast, A₄Z_0.5G hybrid nanocomposite barely revealed any grain boundary cracking besides existence of zirconia nanoparticles (Fig. 5(d)) probably owing to the grain boundary strengthening by the GNP. This corroborates that the zirconia nanoparticles merely contribute in refining the microstructure rather than inducing boundary cracking toughening mechanism in the hybrid nanocomposite. However, grain refinement by zirconia nanoparticles may facilitate the GNP to interact simultaneously with several matrix grains, thus leading to interesting physical interactions, as shown in Fig. 6.

Fig. 6 further shows that physical interactions of the GNP with alumina matrix grains are complicated and interesting toughening mechanism may be anticipated in the hybrid nanocomposites. The sharing and anchoring of several alumina grains by a single GNP segment was observed in the fractured hybrid nanocomposite, as shown in Fig. 6(a). Substantial GNP projection out from the matrix somewhat portrays the existence of pull-out toughing mechanism. Furthermore, the 2D architecture along with large lateral surface facilitated GNP to wrap the matrix grains leading to the GNP locking inside the matrix microstructure, as demonstrated by white arrows in Fig. 6(b). It is assumed that the wrapping of the matrix grains by GNP could lead to additional locking besides the sticking of GNP with the matrix, as shown in Fig. 6(c). These findings elucidate that the interfacial bonding and physical locking between GNP and alumina can synergistically contribute towards activation of pull-out toughening mechanism in the hybrid nanocomposite (A₄Z_0.5G). We further examined the hybrid nanocomposite cross-section using TEM and the results (Fig. 7(a)) manifested rough GNP surface, probably formed due to folding of the GNP segments during hot-pressing. The uneven GNP surface improves the GNP/alumina chemical bonding at the interface. As a consequence, the GNP pull-out from the matrix is extremely difficult due to GNP/alumina larger friction forces. This is also evident from very small GNP pull-out length, as shown in Fig. 6(a) [48]. High-resolution bright-field TEM microstructure of the GNP/alumina interfacial area exhibiting three discrete sections is displayed in Fig. 7(b). In Fig. 7(c), the alumina phase is identifiable from characteristic lattice spacing of 0.26 nm corresponding to the (014) plane, whereas the phase having interplanar spacing of 0.34 nm may be assigned to the (002) planes of the GNP (Fig. 7(b)). An extremely thin interfacial layer at the GNP/alumina junction was also discerned, as marked by a white box in Fig. 7(c). This is in agreement with a report on an intermediate Al₂OC phase formation at the GNP/alumina interfacial region via carbothermal reduction chemistry [49].

In order to gauge the load transfer efficiency of the interfacial area, the GNP/alumina interfacial shear strength was determined by employing a proven mathematical model, as shown in Eq. (4).

$τ_{iss}=E_fe_mβ_{ml}\frac{sinh(\frac{β_{ml}^{x}}{t_{ml}})}{cosh(\frac{β_{ml}^{L}}{2t_{ml}})} $ (4)

$β_{ml}=\sqrt{\frac{2G_{m}}{E_{f}}(\frac{t_{ml}}{T})}$ (5)

$ t_{ml}=\sum_{m=1}^{N}[t+(m-1)ζ]$ (6)

where τiss represents the interfacial shear strength, Ef is the graphene elastic modulus (1 TPa) [49], and e denotes th e alumina matrix with strain (~4%) and β is the effective interfacial stress transfer value [50]. The essential βml and tml parameters were evaluated using Eqs. (5) and (6), in which 2G_m, T, t_ml, ζ and d respectively represent alumina shear modulus (152 GPa), thickness of the matrix alumina, graphene layer thickness (ζ = t + d; t is the graphene wall thickness with a value of 0.32 nm) and the graphene latice spacing (0.34 nm) [50].

From Eq. (4), the GNP/alumina interfacial shear strength value was reckoned to be ~1.2 GPa, which is 30 times higher than the value suggested in case of 1D fibrous reinforcement/alumina ceramic nanocomposites [42]. The high interfacial shear strength may be due to multiple factors such as (i) inherent mechanical properties of GNP, (ii) physical interaction between GNP and alumina matrix grains, and (iii) extremely robust GNP/alumina interface.

For a deeper understanding of the toughening mechanism(s) involved in the hybrid nanocomposite (A₄Z_0.5G), a qualitative schematic model elucidating the GNP contribution towards raising the K_IC values of the hybrid nanocomposite was derived, as illustrated in Fig. 8. As discussed earlier, the grain boundary is the favorable location for the GNP in hybrid nanocomposites microstructure. Also, high lateral surface area and 2D morphology enables GNP to lock numerous matrix grains by wrapping around them. The distinctive GNP/alumina physical interaction could explicitly support GNP to reinforce the alumina matrix grains through anchoring and wrapping so that the GNP could utilize their remarkable flexibility, strength during pull-out and crack-bridging toughening mechanism. Based on both qualitative (GNP/alumina interfacial bonding) and quantitative (high shear strength) approaches, it may be predicted that the nature of pull-out toughening mechanism in hybrid nanocomposites is extremely complicated and quite counterintuitive to those put forth in conventional theories for 1D fibrous and 0D particulate reinforcing materials [51,52].

As depicted in Fig. 8(a), two possible crack propagation scenarios have been proposed based on the microstructural investigation of the hybrid nanocomposites. In the first scenario, the crack is expected to encounter the GNP arranged along the grain boundary (Fig. 8(a-i). The GNP segments obstruct the propagating crack by transferring load to the matrix grains through the thin GNP/alumina interfacial layer and are eventually ruptured (i.e. detachment of individual graphene layers) due to localized stresses of extremely large magnitude, as schematically illustrated in Fig. 8(b)). In the former scenario, GNP elastic deformation is the foremost energy dissipation mechanism prior to the structural failure, as reported elsewhere [53]. The GNP/alumina interfacial layer initially resists crack penetration and the energy dissipation occurs via elastic deformation of the outermost graphene layer in a manner that is quite similar to that explained in case of first crack propagation scenario [54].

In the latter case, the crack may encounter with GNP at the matrix grain boundary with an orientation that is perpendicular to the crack direction, as shown in Fig. 8(a-ii), and therefore, the crack progression proposed in this situation is somewhat more complex than the previous case. When the localized stress further increases to exceed the elastic endurance limit of the outermost graphene layer, the load is transferred to the inner GNP layers upon rupture of the outermost sheets. Since the inner GNP sheets are firmly embedded in the matrix grains via physical anchoring/interlocking, the GNP pullout from the alumina matrix is extremely difficult and unlikely. In the event of the GNP pullout, it initially stretches elastically at lower loading followed by breaking of the C—C bonds that make up the graphene sheets in GNP, at high loads. Eventually, the broken graphene layers slither out from the alumina matrix, as schematically illustrated in Fig. 8(c). However, slithering is a very slow process owing to the reported stick-slip phenomenon between the ruptured sliding graphene layers [22]. In order to accomplish complex stick-slip slithering mechanism, a large amount of energy would be required for stretching the graphene elastically and breaking the strong carbon-carbon bonds as well as overwhelming the friction forces generated between the graphene layers during sliding process. It is deduced from this discussion that, in case of hybrid nanocomposite (A₄Z_0.5G), the zirconia nanoparticles refined the sintered microstructure while supporting the GNP in establishing physical locking with the matrix grains in a 3D configuration (Fig. 8(d)). Together with that, development of a strong GNP/alumina interface contributed towards initiation of slithering mechanism in the hybrid nanocomposite. As a consequence, the efficiency of pullout and crack-bridging toughening mechanisms was greatly enhanced in the hybrid nanocomposites with a remarkable increase in the K_IC value by 155%.

3.5. Analysis of ballistic properties

A ceramic armor system is usually designed after careful, thorough considerations of mechanical properties, lightweight characteristics and ballistic performance of a particular material [30]. The ballistic performance of ceramic-based armor system depends on hardness and fracture toughness and the hybrid nanocomposite sample (A₄Z_0.5G) developed during this work showed noteworthy enhancement in these crucial mechanical properties. The ballistic performance of all the hot pressed samples, in terms ballistic energy dissipation ability (D), was quantitatively measured by taking into consideration the Vickers microhardness value (H_v), modulus of elasticity (E), sonic velocity (c) and fracture toughness (K_IC), in the form of Eq. (7) [55]:

$D=0.36(\frac{H_{v}EC}{K_{IC}^{2}})$(7)

It is evident from Table 2 that the hybrid nanocomposites (A₄Z_0.5G) outperformed all the other samples including monolithic alumina, A₄Z and A_0.5G nanocomposites by exhibiting an increase in the D value by 88%, 71% and 60%, respectively. In ceramic materials, the superior ballistic property is directly related to the lower D-criterion value [55]. It seems that several factors including, homogenous dispersion of zirconia and GNP into alumina matrix, higher densification from hot pressing technique, and grain refinement in the sintered microstructure due to grain boundary pinning by zirconia nanoparticles synergistically contributed to improved ballistic performance of the hybrid nanocomposite (A₄Z_0.5G). It may be deduced that in an event of a bullet impact, the high hardness of the hybrid nanocomposite will blunt the sharp bullet tip, whereas a superior K_IC value will effectively resist the crack propagation through pullout and crack-bridging toughening mechanisms induced by the reinforcing GNP. Therefore, the hot pressed hybrid nanocomposite produced in this study may be a promising contender for armor technology.

4. Conclusions

In this work, hybrid alumina ceramic nanocomposite reinforced with 4 wt% zirconia nanoparticles and 0.5 wt% GNP were fabricated using hot pressing technique. SEM examination revealed that large GNP lateral surface area assisted in attaining uniform dispersion of the zirconia and alumina ceramic nanopowders through an environment-friendly colloidal chemistry route. The homogenously dispersed GNP and zirconia inclusions refined the consolidated microstructure through grain size reduction by 68% via grain pinning effect. The resulting hybrid nanocomposite demonstrated substantial improvement in the K_IC and microhardness values (7.9 MPa m^1/2 and 19.3 GPa) over the monolithic alumina and nanocomposites incorporating either zirconia (A₄Z) or GNP (A_0.5G). Furthermore, the hybrid nanocomposite A₄Z_0.5G showed 88% higher ballistic properties in comparison to the reference alumina sample. The prevalence of high temperature and pressure during hot pressing promoted strong GNP/alumina interfacial bonding and alumina grain boundary anchoring and grain wrapping by the 2D GNP, thus leading to higher interfacial shear strength. The strong GNP/alumina interfacial adhesion and physical locking synergistically enhanced the fracture toughness via GNP pullout and crack-bridging toughness mechanisms.

A schematic model elucidating the reinforcing ability and toughening mechanisms induced by co-addition of zirconia nanoparticles and GNP was proposed based on extensive electron microscopy (both SEM and TEM) of the hybrid nanocomposite. For energy dissipation during crack propagation, the GNP utilize intrinsic elasticity and strength for improved toughness at relatively low applied loads. At much higher localized stresses, however, toughening mechanisms encompass GNP ruptured by means of C—C bonds breaking in the graphene sheets and the subsequent slip-stick atomic movements in the graphene basal plane. It is concluded that the cumulative contributions in terms of alumina matrix grain refinement, the unique alumina/GNP interfacial interaction and mechanical interlocking/anchoring of granular structure by GNP give rise to distinctive toughening mechanisms and improved ballistic properties in the hybrid nanocomposite. The combined higher fracture toughness, hardness, light weight trait and superior ballistic properties can offer the produced hybrid nanocomposites for advanced armor technology and structural applications in the automotive and aircraft industries.

Acknowledgement

The authors would like to extend their sincere appreciation to the Deanship of Scientific Research at King Saud University for funding this research through the Research Group Project No. RGP-283.

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