Epoxy resins were first synthesized in the 1930s, and early applications in coatings, adhesives, and electrical insulation were found. In the 1950s, their use expanded into composite materials, particularly for aerospace and military applications. Since then, the versatility and durability of epoxy have led to its widespread use in various polymer applications, including construction, automotive, and electronics $[1,2]$$[1,2]$[1,2][1,2]$[1,2]$.
Epoxy resins have a variety of inherent qualities due to the highly reactive epoxy groups in their terminal chains. These outstanding qualities make them excellent for use in high-performance applications. Epoxy systems based on diglycidyl ether of bisphenol A (DGEBA) are widely utilized in the plastics sector due to their superior structural qualities, such as high chemical resistance, low shrinkage, and better coating capabilities. [3, 4]. The intrinsic brittleness of thermosetting polymer materials is typically attributed to the high density of cross-linking that occurs during the curing process. Thus, the properties of the cured composite are highly influenced by the type and chemical structure of the monomer and the curing agent used during the curing process. Also, the curing conditions, such as curing time, temperature, and pressure, play a vital role in the synthesized composite [5, 6].
Optimizing the curing process can enhance the final properties of an epoxy system. Lower curing temperatures can result in a thermosetting resin with a lower Tg , as some reactive groups from the epoxy resins or hardeners may not completely react. Post-curing is generally performed at a temperature higher than the initial curing temperature to achieve optimal crosslinking. This process leads to a resin with superior mechanical properties and low shrinkage, which improves stability. Additionally, the resultant resins can offer better resistance in surface coatings and better adhesion strength making them ideal for various industrial applications, including paints, adhesives, and high-performance membranes [7, 8]. It is also known that polymeric-based nanocomposite materials cannot be used for high-performance applications because of their limited properties. This limitation can be overcome by introducing organic/inorganic nanoparticles. GNPs offer significant potential for multifunctional properties. These GNPs are made up of graphene sheet stacks of platelet structure, similar to the planar form found in carbon nanotubes. GNPs can improve the mechanical properties of matrix materials, including stiffness, strength, and surface hardness [9]. h-BNs have tremendous potential for multifunctional applications. h-BN is constructed of minute hexagonal layers that are physically similar to graphene but contain alternating boron and nitrogen atoms. This unusual structure gives $\mathrm{h}-\mathrm{BN}$$\mathrm{h}-\mathrm{BN}$h-BN\mathrm{h}-\mathrm{BN}$\mathrm{h}-\mathrm{BN}$ outstanding thermal and electrical insulating qualities. Furthermore, $\mathrm{h}-\mathrm{BN}$$\mathrm{h}-\mathrm{BN}$h-BN\mathrm{h}-\mathrm{BN}$\mathrm{h}-\mathrm{BN}$ nanofillers can improve the mechanical properties of the matrix material by enhancing heat conductivity, lubricity, and resistance to wear and corrosion [10]. The ultrasonication process can be used to achieve an optimal distribution of nanoparticles (NPs) within the base material, ensuring the maximum dispersion of NPs in the holding matrix [11]. FEM
analysis with Ansys software is critical for designing polymer nanocomposites because it allows simulation of their mechanical and thermal properties at the microstructural level. Ansys optimizes material composition, predicts performance, and identifies probable failure areas by simulating the interactions between the polymer matrix and nanofillers [15,23,24]. Olszowska et al. found that composites containing GNPs and micro-sized carbon foam in an epoxy matrix improved thermal conductivity by $20\mathrm{\%}$$20\mathrm{\%}$20%20 \%$20\mathrm{\%}$, mechanical strength by $15\mathrm{\%}$$15\mathrm{\%}$15%15 \%$15\mathrm{\%}$, and friction coefficient by $30\mathrm{\%}$$30\mathrm{\%}$30%30 \%$30\mathrm{\%}$ [9]. Choukimath et al. reported an increase of $49.25\mathrm{\%}$$49.25\mathrm{\%}$49.25%49.25 \%$49.25\mathrm{\%}$ in the tensile strength of h-BN based epoxy nanocomposites prepared with varying compositions of h-BN from 0.1 to $0.5\mathrm{wt}\mathrm{\%}$$0.5\mathrm{wt}\mathrm{\%}$0.5wt%0.5 \mathrm{wt} \%$0.5\mathrm{wt}\mathrm{\%}$ [12]. Jahani et al. examined the effect of post-curing on the mechanical properties of a structural epoxy glue at various temperatures. Their work concluded that curing and post-curing below the adhesive's Tg improve mechanical characteristics, whereas temperatures over Tg cause degradation. Testing temperatures above ${20}^{\circ }\mathrm{C}$${20}^{\circ }\mathrm{C}$20^(@)C20^{\circ} \mathrm{C}${20}^{\circ }\mathrm{C}$ have a negative impact on tensile and compressive characteristics, with considerable reductions reported at temperatures near and beyond Tg [13]. Min, C. et al. describe the creation of a three-dimensional interconnected graphene architecture (3DGA) reinforced epoxy composite. Using resin transfer molding and an in-situ ultrasonic technique, the 3DGA was evenly distributed within the epoxy matrix, resulting in improved thermal stability, mechanical strength, and tribological properties. Compared to typical graphene oxide composites, the 3DGA/EP composite showed a ${27.1}^{\circ }\mathrm{C}$${27.1}^{\circ }\mathrm{C}$27.1^(@)C27.1^{\circ} \mathrm{C}${27.1}^{\circ }\mathrm{C}$ rise in glass transition temperature, a $41.1\mathrm{\%}$$41.1\mathrm{\%}$41.1%41.1 \%$41.1\mathrm{\%}$ increase in tensile strength, and a $38.1\mathrm{\%}$$38.1\mathrm{\%}$38.1%38.1 \%$38.1\mathrm{\%}$ decrease in wear rate, indicating its potential for highperformance applications [14]. Choukimath et al. reported a significant increase in tensile, flexural load bearing, and thermal stability of GNP and h-BN embedded epoxy nanocomposites. The inclusion of GNPs increased the load-bearing capacity by $265\mathrm{\%}$$265\mathrm{\%}$265%265 \%$265\mathrm{\%}$ and the flexural strength by $165\mathrm{\%}$$165\mathrm{\%}$165%165 \%$165\mathrm{\%}$ at a concentration of $0.2\mathrm{wt}\mathrm{\%}$$0.2\mathrm{wt}\mathrm{\%}$0.2wt%0.2 \mathrm{wt} \%$0.2\mathrm{wt}\mathrm{\%}$. When coupled with $\mathrm{h}-\mathrm{BN}$$\mathrm{h}-\mathrm{BN}$h-BN\mathrm{h}-\mathrm{BN}$\mathrm{h}-\mathrm{BN}$ at $0.6\mathrm{wt}\mathrm{\%}$$0.6\mathrm{wt}\mathrm{\%}$0.6wt%0.6 \mathrm{wt} \%$0.6\mathrm{wt}\mathrm{\%}$, the nanocomposite exhibited increased thermal stability, withstanding higher temperatures as confirmed by thermogravimetric analysis. These enhancements make the nanocomposite suitable for high-temperature applications, including compressor blades in gas turbine engines[15].
In this study, a comprehensive investigation is conducted on hybrid epoxy nanocomposites reinforced with GNP) and hBN nanofillers, which are subjected to various post-curing temperatures. Unlike prior studies, the synergistic effects of dual nanofiller reinforcement and different post-curing temperatures on the mechanical, thermal, and fracture properties of the composites are examined. Additionally, experimental data combined with ANSYS simulation provides strong confirmation of the observed behaviors, revealing new understandings for enhancing material properties in high-performance applications. Subsequent sections provide information on materials, methodology, specimen preparation, experimental details, test procedure and analysis.

The epoxy matrix material (branded as Lapox-L12 and paired with K6 hardener) used in this study was sourced from Atul Private Limited's Polymer division Gujarat, India. The GNPs and h-BNs used in the study had a minimum purity of $99.5\mathrm{\%}$$99.5\mathrm{\%}$99.5%99.5 \%$99.5\mathrm{\%}$ and $99.9\mathrm{\%}$$99.9\mathrm{\%}$99.9%99.9 \%$99.9\mathrm{\%}$ respectively. The properties of GNP and h-BN used are displayed in Tab. 1.

The nano-composites were synthesized using an in-situ polymerization process. Three types of epoxy-based nanocomposites were prepared, one with GNPs, another with h-BNs and the third with a combination of both GNPs and h-BNs. The nanocomposite composition is displayed in Tab. 2. After adding the appropriate amount of nanofillers and epoxy matrix by weight percentage, the mixture was subjected to sonication for 60 minutes to improve the nanofillers dispersion in the matrix [15]. Further, $10\mathrm{\%}$$10\mathrm{\%}$10%10 \%$10\mathrm{\%}$ by weight of aliphatic polyamine hardener (K6) was added to the epoxy base to initiate the polymerization process. This final mixture was cast into the aluminum mold of size $230\mathrm{mm}\times 164\mathrm{mm}\times 3\mathrm{mm}$$230\mathrm{mm}\times 164\mathrm{mm}\times 3\mathrm{mm}$230mmxx164mmxx3mm230 \mathrm{~mm} \times 164 \mathrm{~mm} \times 3 \mathrm{~mm}$230\mathrm{mm}\times 164\mathrm{mm}\times 3\mathrm{mm}$. These cast plates were allowed to cure for 24 hours. After the curing process, the composite plates were removed from the molds and were cut to the ASTM specifications for different test procedures. Further, these prepared composites were subjected to post-curing in a hot air oven for 120 minutes at different post-curing temperatures $({80}^{\circ }\mathrm{C}$$\left({80}^{\circ }\mathrm{C}\right.$(80^(@)C:}\left(80^{\circ} \mathrm{C}\right.$({80}^{\circ }\mathrm{C}$, ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ and ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$ ). Details of the post-curing process are displayed in Tab. 3.

The synthesized nanocomposites were evaluated using a range of characterization techniques. The experimental details are as follows.

The specifications of the test specimens for tensile, flexure impact and fracture tests are depicted in Fig. 1. For all mechanical and fracture tests, 3 samples were tested and an average with standard error was used in plotting the relevant graphs.

Finite Element (FE) analysis was used to estimate mechanical properties, which were then validated using experimental test data. The simulation studies were carried out utilizing the ANSYS 2021 R2 software. Nanocomposites were designed using ANSYS Workbench's material designer module by specifying essential parameters such as Young's modulus, Poisson's ratio, and ultimate strength. Tensile and flexural models were meshed using solid 186 elements, which were simulated followed by the application of boundary conditions (BC) [23,24].

The Raman spectroscopy results indicate the presence of epoxide functional groups in the epoxy material under investigation. The characteristic Raman peaks observed at specific wavenumbers, consistent with vibrational modes associated with epoxide moieties, serve as a distinctive signature. These findings provide crucial insights into the molecular composition of the epoxy, confirming the presence of epoxide groups and nanoparticles (GNPs and h-BNs) in the prepared nanocomposites. Fig. 2 shows the Raman spectra of PE, GNP, h-BN and GH-based nanocomposites.
In Raman spectroscopy, the presence of epoxide (or epoxy) functional groups in a material is often characterized by specific vibrational modes associated with the C-O-C bonds in the epoxy ring. C-O-C stretching was observed at $1280-1230{\mathrm{cm}}^{-1}$$1280-1230{\mathrm{cm}}^{-1}$1280-1230cm^(-1)1280-1230 \mathrm{~cm}^{-1}$1280-1230{\mathrm{cm}}^{-1}$ while $\mathrm{C}-\mathrm{O}-\mathrm{C}$$\mathrm{C}-\mathrm{O}-\mathrm{C}$C-O-C\mathrm{C}-\mathrm{O}-\mathrm{C}$\mathrm{C}-\mathrm{O}-\mathrm{C}$ bending was observed at $800-900{\mathrm{cm}}^{-1}$$800-900{\mathrm{cm}}^{-1}$800-900cm^(-1)800-900 \mathrm{~cm}^{-1}$800-900{\mathrm{cm}}^{-1}$ thus confirming the epoxide presence in the base material [20]. The presence of GNP is witnessed near $1580{\mathrm{cm}}^{-1}$$1580{\mathrm{cm}}^{-1}$1580cm^(-1)1580 \mathrm{~cm}^{-1}$1580{\mathrm{cm}}^{-1}$ representing the graphitic sp2 carbon-carbon stretching in the G Band followed by the presence of the D band at $1350{\mathrm{cm}}^{-1}$$1350{\mathrm{cm}}^{-1}$1350cm^(-1)1350 \mathrm{~cm}^{-1}$1350{\mathrm{cm}}^{-1}$ and 2 D band at $2700{\mathrm{cm}}^{-1}$$2700{\mathrm{cm}}^{-1}$2700cm^(-1)2700 \mathrm{~cm}^{-1}$2700{\mathrm{cm}}^{-1}$ [21]. Further, the presence of $\mathrm{h}-\mathrm{BNs}$$\mathrm{h}-\mathrm{BNs}$h-BNs\mathrm{h}-\mathrm{BNs}$\mathrm{h}-\mathrm{BNs}$ in the prepared composites was witnessed at $1365{\mathrm{cm}}^{-1}$$1365{\mathrm{cm}}^{-1}$1365cm^(-1)1365 \mathrm{~cm}^{-1}$1365{\mathrm{cm}}^{-1}$ of E 2 g In-Plane Mode vibration and $800{\mathrm{cm}}^{-1}$$800{\mathrm{cm}}^{-1}$800cm^(-1)800 \mathrm{~cm}^{-1}$800{\mathrm{cm}}^{-1}$ wavenumbers of A1g Out-of-Plane Mode, which corresponds to the vibration of boron and nitrogen in a hexagonal lattice [22].

FTIR spectroscopy was also used to evaluate the functional groups of epoxy and nanofillers used in the studies. Fig. 3 shows the FT-IR spectra of PE, GNP, h-BN and GH-based nanocomposites. A characteristic peak for the C-O-C ring stretching vibration in the epoxy group was observed in the range of $850-1000{\mathrm{cm}}^{-1}$$850-1000{\mathrm{cm}}^{-1}$850-1000cm^(-1)850-1000 \mathrm{~cm}^{-1}$850-1000{\mathrm{cm}}^{-1}$ and $\mathrm{C}-\mathrm{O}-\mathrm{C}$$\mathrm{C}-\mathrm{O}-\mathrm{C}$C-O-C\mathrm{C}-\mathrm{O}-\mathrm{C}$\mathrm{C}-\mathrm{O}-\mathrm{C}$ Ring Bending was observed around $430-500{\mathrm{cm}}^{-1}$$430-500{\mathrm{cm}}^{-1}$430-500cm^(-1)430-500 \mathrm{~cm}^{-1}$430-500{\mathrm{cm}}^{-1}$ confirming the presence of the epoxy functional group. The presence of GNP in the composite is witnessed by the presence of G band at $1580{\mathrm{cm}}^{-1}$$1580{\mathrm{cm}}^{-1}$1580cm^(-1)1580 \mathrm{~cm}^{-1}$1580{\mathrm{cm}}^{-1}$ which corresponds to $\mathrm{C}-\mathrm{C}$$\mathrm{C}-\mathrm{C}$C-C\mathrm{C}-\mathrm{C}$\mathrm{C}-\mathrm{C}$ stretching in graphene. The peaks of (A1g Out-of-Plane Mode) at $800{\mathrm{cm}}^{-1}$$800{\mathrm{cm}}^{-1}$800cm^(-1)800 \mathrm{~cm}^{-1}$800{\mathrm{cm}}^{-1}$ correspond to the presence of $\mathrm{h}-\mathrm{BN}$$\mathrm{h}-\mathrm{BN}$h-BN\mathrm{h}-\mathrm{BN}$\mathrm{h}-\mathrm{BN}$ in the composite material [22].

DSC analysis was performed on select cured and post-cured specimens. The selection of samples for DSC analysis was based on the results from tensile tests, where samples (HBN3, GNP2 and GH3) post cured at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ exhibited highest
tensile strength in their corresponding configurations. All specimens that were subjected to analysis underwent thermal exposure ranging from ambient to ${400}^{\circ }\mathrm{C}$${400}^{\circ }\mathrm{C}$400^(@)C400^{\circ} \mathrm{C}${400}^{\circ }\mathrm{C}$ in a nitrogen environment. Fig. 4 (a) shows the DSC thermographs of cured composites and Fig. 4 (b) shows the DSC thermographs subjected to post-curing ( ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ ).
The DSC curves show no substantial shift in Tg for room-temperature cured composites (Fig. 4 (a)), as evident from the thermograph the Tg valves of cured composites lay in the range of ${63}^{\circ }\mathrm{C}$${63}^{\circ }\mathrm{C}$63^(@)C63^{\circ} \mathrm{C}${63}^{\circ }\mathrm{C}$ to ${66}^{\circ }\mathrm{C}$${66}^{\circ }\mathrm{C}$66^(@)C66^{\circ} \mathrm{C}${66}^{\circ }\mathrm{C}$. The prominent exothermic peaks in the cured specimens have diminished in the post-cured specimens, indicating that post-curing may have completed the reaction or altered the thermal properties. From Fig. 4 (b), the Tg of HBN3, GNP2 and GH3 at ${112.85}^{\circ }\mathrm{C},{115.45}^{\circ }\mathrm{C}$${112.85}^{\circ }\mathrm{C},{115.45}^{\circ }\mathrm{C}$112.85^(@)C,115.45^(@)C112.85^{\circ} \mathrm{C}, 115.45^{\circ} \mathrm{C}${112.85}^{\circ }\mathrm{C},{115.45}^{\circ }\mathrm{C}$ and ${110.2}^{\circ }\mathrm{C}$${110.2}^{\circ }\mathrm{C}$110.2^(@)C110.2^{\circ} \mathrm{C}${110.2}^{\circ }\mathrm{C}$ suggest specific transitions or reactions characteristic of these formulations. The curve shapes indicate that post-curing may have enhanced the thermal stability of the materials since there are fewer variations in heat flow compared to the cured specimens. The average Tg value of the post-cured specimens increases by $174\mathrm{\%}$$174\mathrm{\%}$174%174 \%$174\mathrm{\%}$ compared to cured composites. This shift in the DSC curves after post-curing could be due to an increase in the cross-link density or a change in the chemical structure of the composites, leading to different thermal properties. The difference in the DSC curves before and after post-curing can also provide insights into the curing process's kinetics and the post-curing treatment's efficiency [5, 6]. The PE of postcured specimen shows no significant improvement in the Tg value compared to PE of the cured composite. The thermographs suggest that the addition of GNP and h-BN has a synergistic effect on the thermal behavior of the epoxy composites [26].

This study examines the relationship between reinforcements and the matrix material ultimate stress of PE and various reinforced epoxy specimens subjected to tensile tests after post-curing at three different temperatures: ${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$80^(@)C,120^(@)C80^{\circ} \mathrm{C}, 120^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$, and ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$. Fig. 5 shows the variation of ultimate stress in all specimens.
At ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ post curing PE exhibits tensile stress of 30.7 MPa , GNP1, HBN3 and GH3 showed an increase in stress-bearing capacity by $174\mathrm{\%}(53.6\mathrm{MPa}),135\mathrm{\%}(41.5\mathrm{MPa})$$174\mathrm{\%}(53.6\mathrm{MPa}),135\mathrm{\%}(41.5\mathrm{MPa})$174%(53.6MPa),135%(41.5MPa)174 \%(53.6 \mathrm{MPa}), 135 \%(41.5 \mathrm{MPa})$174\mathrm{\%}(53.6\mathrm{MPa}),135\mathrm{\%}(41.5\mathrm{MPa})$ and $157\mathrm{\%}(48.4\mathrm{MPa})$$157\mathrm{\%}(48.4\mathrm{MPa})$157%(48.4MPa)157 \%(48.4 \mathrm{MPa})$157\mathrm{\%}(48.4\mathrm{MPa})$ respectively as compared to PE. This improvement is due to the effective load transmission and strengthening provided by GNPs, HBNs and GH (GNPs $+\mathrm{h}-\mathrm{BNs}$$+\mathrm{h}-\mathrm{BNs}$+h-BNs+\mathrm{h}-\mathrm{BNs}$+\mathrm{h}-\mathrm{BNs}$ ) nanofillers within the epoxy matrix respectively. Specimens post-cured at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ have fared better tensile load-bearing capacity when compared to PE ( 40 MPa ). This could be due to the increase in Tg of the specimens. Higher Tg increases crosslinking, chain stiffness, and strong intermolecular interactions resulting in better stress-bearing capacity before failure [15]. GNP1 and HBN3 showed an increase in stress-bearing capacity by $181.5\mathrm{\%}(72.6\mathrm{MPa})$$181.5\mathrm{\%}(72.6\mathrm{MPa})$181.5%(72.6MPa)181.5 \%(72.6 \mathrm{MPa})$181.5\mathrm{\%}(72.6\mathrm{MPa})$ and $141.0\mathrm{\%}(56.4\mathrm{MPa})$$141.0\mathrm{\%}(56.4\mathrm{MPa})$141.0%(56.4MPa)141.0 \%(56.4 \mathrm{MPa})$141.0\mathrm{\%}(56.4\mathrm{MPa})$ respectively compared to PE. At ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$ post curing, a considerable reduction in ultimate stress is observed; this significant drop is due to the
thermal instability of the material when exposed to higher post-curing temperature than the Tg limit of the base material (epoxy). However, GNP1 and HBN3 showed an increase in stress-bearing capacity by $202.4\mathrm{\%}$$202.4\mathrm{\%}$202.4%202.4 \%$202.4\mathrm{\%}$ ( 50.2 MPa) and $210.0\mathrm{\%}$$210.0\mathrm{\%}$210.0%210.0 \%$210.0\mathrm{\%}$ ( 52.1 MPa ) respectively compared to PE ( 24.8 MPa ). However, GH1 showed a decrease in stress-bearing capacity by $22.8\mathrm{\%}$$22.8\mathrm{\%}$22.8%22.8 \%$22.8\mathrm{\%}$ (when compared to PE at ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ ) due to thermal softening at higher post-curing temperatures [5].

To analyze the surface morphology, dispersion of filler, and their interaction with the matrix, the fractured tensile specimens were examined using SEM. The micrographs of the pure epoxy and one specimen from each nanocomposite configuration were studied and shown in Fig. 6. The specimens exhibiting higher tensile strength among their corresponding filler configuration were selected. From Fig. 6 (a), the pure epoxy surface showed a smooth surface that implies a typical brittle failure. The samples loaded with filers (Fig. 6 (b) - Fig. 6 (c)) exhibited rough surfaces and the cleavage planes were observed. The morphology of GNP nanocomposite (Fig. 6 (b)) exhibits a very rough surface with several cleavage planes, representing regions capable of absorbing fracture energy that results in a higher resistance for crack propagation. The surface of $h-BN$$h-BN$h-BNh-B N$h-BN$ nanocomposites (Fig. 6 (c)) is relatively smooth and exhibits cleavage planes. The hybrid nanocomposite (Fig. 6 (d)) surface exhibits rough surfaces in some regions and relatively smooth surfaces with cleavage planes in both regions. The higher strength of GNP nanocomposites compared to h-BN and hybrid nanocomposites can be attributed to the higher strength and modulus of GNPs in comparison with h-BN particles [25, 26].

This study examines the relationship between reinforcements and the matrix material flexural stress of PE and various reinforced epoxy specimens subjected to flexure tests after post-curing at three different temperatures: ${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$80^(@)C,120^(@)C80^{\circ} \mathrm{C}, 120^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$, and ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$. Fig. 7 shows the variation of flexural stress in all specimens.

At ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ post-curing PE exhibits a flexural stress of 40.63 MPa , GNP2 shows a substantial increase in flexural stress of 139.58 MPa , an increase of $243\mathrm{\%}$$243\mathrm{\%}$243%243 \%$243\mathrm{\%}$ compared to PE. This improvement is due to the effective load transmission and strengthening of GNP nanofillers within the epoxy matrix. Specimens post-cured @ ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ have fared better tensile loadbearing capacity when compared with ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ (barring GNP2) post-cured specimens and PE, this could be due to the increase in Tg of the specimens. Higher Tg increases crosslinking, chain stiffness, and strong intermolecular interactions resulting in better stress-bearing capacity before failure [15]. GNP2, GNP3 and HBN2 showed an increase in stress-bearing capacity by $256\mathrm{\%},240\mathrm{\%}$$256\mathrm{\%},240\mathrm{\%}$256%,240%256 \%, 240 \%$256\mathrm{\%},240\mathrm{\%}$ and $234\mathrm{\%}$$234\mathrm{\%}$234%234 \%$234\mathrm{\%}$ respectively when compared to PE ( 33.33 MPa ). Specimens post-cured at ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$, drop in stressbearing capacity was observed. This is due to the thermal instability of the material when exposed to higher post-curing temperatures than the Tg limit of the base material (epoxy). However, GNP2 and HBN2 showed an increase in stressbearing capacity by $298\mathrm{\%}$$298\mathrm{\%}$298%298 \%$298\mathrm{\%}$ and $201\mathrm{\%}$$201\mathrm{\%}$201%201 \%$201\mathrm{\%}$ respectively compared to PE ( 28.54 MPa ) but showed a reduction in bearing strength by $18.6\mathrm{\%}$$18.6\mathrm{\%}$18.6%18.6 \%$18.6\mathrm{\%}$ and $38.4\mathrm{\%}$$38.4\mathrm{\%}$38.4%38.4 \%$38.4\mathrm{\%}$ when compared with GNP2 (of ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ ). GH1 and GH3 showed a decrease in stress-bearing capacity by $11.5\mathrm{\%}$$11.5\mathrm{\%}$11.5%11.5 \%$11.5\mathrm{\%}$ and $15.4\mathrm{\%}$$15.4\mathrm{\%}$15.4%15.4 \%$15.4\mathrm{\%}$ respectively when compared to PE (of ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ ). This decrease is attributed to thermal degradation resulting in a drop in flexural performance [5,6].

This study examines the relationship between reinforcements and the matrix material impact strength of PE and various reinforced epoxy specimens subjected to impact tests after post-curing at three different temperatures: ${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$80^(@)C,120^(@)C80^{\circ} \mathrm{C}, 120^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$, and ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$. GNP and h-BN based composites outperformed in absorbing the impact before the break point compared to the PE. This is mainly due to the proper dispersion of NPs in the holding matrix (epoxy) ensuring a stronger interface between them [26]. Fig. 8 shows the impact resistance in terms of impact strength against all specimens subjected to post-curing temperatures.

At ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ post curing all specimens (except GH1) exhibited better resistance to impact loading compared to PE ( $72.32\mathrm{J}/\mathrm{m}$$72.32\mathrm{J}/\mathrm{m}$72.32J//m72.32 \mathrm{~J} / \mathrm{m}$72.32\mathrm{J}/\mathrm{m}$ ). This could be related to the material's thermal deterioration limitations, which affect its structural integrity. GNP1 showed an increase in $312\mathrm{\%}$$312\mathrm{\%}$312%312 \%$312\mathrm{\%}$ impact resistance when compared with PE. This behavior could be due to efficient interfacial behavior between the GNP and the matrix material. All specimens post-cured at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ have exhibited good resistance before failure. HBN1, HBN2 and HBN3 outperformed the impact resistance when compared with GNP reinforced composites, this might be due to the uniform distribution of h-BNs in the matrix material as well as higher glass transition temperature or a composition that resists thermal softening better. Specimens post-cured at ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$, showed a significant decrease in impact
strength, with only a few specimens maintaining values slightly above that of PE post-cured at ${80}^{\circ }\mathrm{C}(72.32\mathrm{J}/\mathrm{m})$${80}^{\circ }\mathrm{C}(72.32\mathrm{J}/\mathrm{m})$80^(@)C(72.32J//m)80^{\circ} \mathrm{C}(72.32 \mathrm{~J} / \mathrm{m})${80}^{\circ }\mathrm{C}(72.32\mathrm{J}/\mathrm{m})$. This drop in impact resistance suggests that most test specimens have had higher crosslinking reactions leading to brittle nature [26].

The fracture test investigates the relationship between reinforcements and the matrix material fracture strength of PE and various reinforced epoxy specimens subjected to post-curing at three different temperatures: ${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$80^(@)C,120^(@)C80^{\circ} \mathrm{C}, 120^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C},{120}^{\circ }\mathrm{C}$, and ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$. Fig. 9 shows the fracture toughness against all specimens subjected to post-curing temperatures.
The fracture toughness of PE specimen post-cured at ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ is $2.86\mathrm{MPa}.{\mathrm{m}}^{1/2}$$2.86\mathrm{MPa}.{\mathrm{m}}^{1/2}$2.86MPa.m^(1//2)2.86 \mathrm{MPa} . \mathrm{m}^{1 / 2}$2.86\mathrm{MPa}.{\mathrm{m}}^{1/2}$ and is used as a reference for comparison. GNP reinforced composites perform well due to strong bonding with epoxy matrix. GNP2 showed an increase in fracture toughness by $60\mathrm{\%}$$60\mathrm{\%}$60%60 \%$60\mathrm{\%}$ when compared with PE. Specimens post-cured at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$, showed further improvement in fracture strength, particularly in h-BN reinforced samples. This improvement is because of the uniform dispersion of h-BN and its higher glass transition temperature, enhancing fracture resistance. However, specimens post-cured at ${160}^{\circ }\mathrm{C}$${160}^{\circ }\mathrm{C}$160^(@)C160^{\circ} \mathrm{C}${160}^{\circ }\mathrm{C}$, exhibited a significant reduction in fracture strength, indicating thermal degradation. All nanocomposites (Except GH3) post-cured at 160 C exhibited an increase in fracture toughness compared to PE [13, 24].

In the material designer module of Ansys Workbench, a new material was developed in the engineering data section, as illustrated in Fig. 10 (a). Basic properties, such as Young's modulus, Poisson's ratio, and ultimate strength, derived from tensile tests and flexural tests were employed to create the matrix material [23]. Additional elasticity properties, including bulk modulus and shear modulus, were determined using the built-in methods provided by the material designer. Representative volume elements (RVEs) were generated for each GNP1, HBN3, and GH3 of tensile models and GNP2, HBN2, and GH2 of flexural models. The mechanical properties of epoxy, GNP, and h-BN, detailed in Tab. 1, were utilized to construct the RVEs, the meshing of RVEs is depicted in Fig. 10 (b). This material was used as input for structural analysis in ANSYS Workbench. 3D models for tensile and flexural test analysis were created in SolidWorks using the specifications provided in Fig. 1 [24].

The tensile and flexural models were meshed with Solid 186 element as shown in Figs. 11 (a) & 11 (b) respectively. The BC of tensile tests was done with one end clamped, and axial force was applied on the other end, as shown in Fig. 12 (a). The

BC of flexural tests was done with two ends clamped and an axial force was applied in the mid-span of the test specimen as shown in Fig. 12 (b) [23, 24].

Figs. 13 (a), 14 (a), and 15 (a) represent the simulated tensile test results of GNP1, HBN3 and GH3 specimens respectively. Figs. 13 (b), 14 (b) and 15 (b) represent the simulated flexural results of GNP2, HBN2 and GH2 specimens respectively. Tab. 4 illustrates the simulated $\mathrm{v}/\mathrm{s}$$\mathrm{v}/\mathrm{s}$v//s\mathrm{v} / \mathrm{s}$\mathrm{v}/\mathrm{s}$ experimental results of the tensile and flexural models. Overall, the variations in both tensile and flexural tests between experimental and simulated results for all nanocomposites remained within the acceptable range of $20\mathrm{\%}$$20\mathrm{\%}$20%20 \%$20\mathrm{\%}$ [15].

The study confirms the incorporation of GNP and h-BN nanofillers in an epoxy matrix improves their mechanical properties, such as tensile, flexural, and impact strength. This enhancement is due to increased interfacial adhesion between the nanofillers and the epoxy matrix. The epoxy composites infused with varied concentrations of GNPs
and h-BN were evaluated, demonstrating that changing the nanofiller content can successfully customize the composite's mechanical, fracture, and thermal properties to match specific application requirements.
The multi-filler system, which included GNP and h-BN, outperformed PE in terms of mechanical, fracture, and thermal properties. During tensile loading, GNP1 at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ absorbed $136\mathrm{\%}$$136\mathrm{\%}$136%136 \%$136\mathrm{\%}$ more stress before failing, and GNP2 at ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ absorbed $243\mathrm{\%}$$243\mathrm{\%}$243%243 \%$243\mathrm{\%}$ more flexural loading. Similarly, GNP1 at ${80}^{\circ }\mathrm{C}$${80}^{\circ }\mathrm{C}$80^(@)C80^{\circ} \mathrm{C}${80}^{\circ }\mathrm{C}$ required $212\mathrm{\%}$$212\mathrm{\%}$212%212 \%$212\mathrm{\%}$ greater impact stress before failing, whereas GNP2 at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ showed a $71\mathrm{\%}$$71\mathrm{\%}$71%71 \%$71\mathrm{\%}$ increase in resistance to failure. FTIR, SEM and Raman spectroscopy were used to confirm the presence and dispersion of nanofillers in the epoxy matrix, providing more information about the composite structure and filler-matrix interactions.
It has also been found that post-curing temperatures substantially impact the composite's mechanical and thermal properties. Specimens post-cured at ${120}^{\circ }\mathrm{C}$${120}^{\circ }\mathrm{C}$120^(@)C120^{\circ} \mathrm{C}${120}^{\circ }\mathrm{C}$ showed improved mechanical and thermal performance, particularly in Tg and tensile strength. These improved nanocomposite materials indicate a considerable promise for high-performance applications in aerospace and automotive space where durability and thermal resistance are crucial. Furthermore, simulation analysis using ANSYS software showed a high match with experimental results, with variances falling within a $20\mathrm{\%}$$20\mathrm{\%}$20%20 \%$20\mathrm{\%}$ range.

The authors gratefully acknowledge the organization for providing the essential facilities at the Center for Materials Science and the School of Mechanical Engineering, KLE Technological University, Hubballi, Karnataka, India.

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