INTRODUCTIONThe rapid development of microelectronic and optoelectronic devices has prompted the end products to gradually move towards high integration, miniaturization, high-speed and high-frequency.[1−3] Owing to the increasing power density per unit volume, a large amount of heat would accumulate readily during service of the devices. Thereby, high requirements for excellent heat dissipation capability and heat resistance,[4,5] have been put forwards for the printed circuit boards (PCBs) that act as the carrier of electronic elements.Glass fiber reinforced epoxy (GFREP) composites are widely applied as the organic substrate of PCBs, on account of their superior electrical insulation, good mechanical properties, low cost, and easy processability. However, the inherent low thermal conductivity (0.2−0.3 W·m−1·K−1) and low Tg (130−150 °C) are unable to meet the increasing demands. Besides, the thermal stress and external impact often cause internal cracks and interfacial debonding in the composite laminates,[6,7] while the thermosetting feature of the permanent cross-linked epoxy networks hinder their intrinsic self-healing capability. Hence, imparting GFREP composite with high thermal conductivity, heat resistance and healability is highly desired to promote the upgradation of comprehensive performance.Generally, the heat dissipation issue can be tackled via the addition of thermally conductive ceramic fillers (e.g., Al2O3 and boron nitride (BN)) in the composites.[8−18] In order to form perfect heat transport networks, loading of the inorganic particles has to be higher than 50 wt% in most cases. Under the circumstances, the mechanical strength and electrical insulation properties have to be significantly decayed.[19]On the other hand, thermally conductive polymer composites containing dynamic reversible bonds are allowed to rearrange the topological networks of the matrices by the exchange reactions of built-in reversible bonds under external stimuli,[20] and drive recontact of thermally conductive particles located on the fracture surfaces.[6,21] As a result, mechanical properties and heat conduction paths of the composites are allowed to be synchronously restored. So far, most of the reported achievements in the area suffer from insufficient mechanical strength (50 MPa) and low Tg (30 °C) due to the presence of reversible bonds, which greatly limit their application in the high-end products.[22−30] To maintain material stability of intrinsic self-healing polymers bearing dynamic reversible bonds in service and promote their applications, the thermally triggered reversible bonds with high universality and reversible temperature exceeding room temperature are desirable. Comparatively, the β-hydroxyl ester bonds can be in situ generated during the curing reactions between epoxies and anhydride/carboxylic acid hardeners, which excludes the complicated synthesis of specific monomers and have been widely applied in the design of healable polymers.[21,31]Herein, a thermally conductive, healable GFREP composite with enhanced Tg and thermal conductivity is fabricated by combining the dynamic crosslinked epoxy carrying reversible β-hydroxyl ester bonds (Fig. 1a), glass fiber cloth and nanometer/micrometer sized modified BN platelets. Firstly, the glass fiber cloth (GFC) and alkali activated boron nitride nanosheets (BNNSs) are treated by the negatively charged poly(acrylic acid) (PAA) and positively charged polyethyleneimine (PEI), respectively, followed by the electrostatic self-assembly to yield the GFC@BNNSs with well-designed in-plane heat conduction paths, so as to diminish the interfacial heat resistance (Fig. 1b). Afterwards, the GFC@BNNSs are impregnated by dip coating solution consisting of lower concentration (3-glycidyloxypropyl) trimethoxysilane grafted hexagonal boron nitride micron sheets (h-BN@KH560), 3,4-epoxycyclohexylmethyl 3,4-epoxycyclohexanecarboxylate (CY-179) and hexahydro-4-methylphthalic anhydride (MHHPA), which helps to construct the through-plane heat transport pathways and avoids excessive deterioration of mechanical performances.Fig 1(a) Synthesis mechanism of dynamic epoxy network; (b) Fabrication of the healable thermally conductive GFREP; (c) Transesterification reaction induced network rearrangement.The rigid cycloaliphatic skeleton of CY-179 and MHHPA are expected to improve the heat tolerance of the resultant dynamic epoxy networks and composites. The glass fiber cloth adsorbed with nano-boron nitride can enhance mechanical strength and thermal conductivity of the composite using relatively fewer boron nitride fillers. Meantime, the low viscosity of CY-179 and MHHPA would bring benefit to the impregnation and lamination of GFC manufacturing. Besides, the epoxy moieties of h-BN@KH560 can participate in the curing reaction, ensuring strong filler-matrix interfacial interaction. After being overlapped and hot-pressed to undergo epoxy-anhydride crosslinking reaction (Fig. 1a), the healable GFREP composite with high in-plane (k∥) and through-plane (k⊥) thermal conductivity is produced at a proper BN content. The reversible transesterification reaction induced topological network rearrangement (Fig. 1c) imparts crack healability to the GFREP, allowing the restoration of mechanical and thermally conductive properties after internal injure.EXPERIMENTALMaterials3,4-Epoxycyclohexylmethyl 3,4-epoxycyclohexanecarboxylate (CY-179), hexahydro-4-methylphthalic anhydride (MHHPA), zinc acetylacetonate hydrate (ZAH) were obtained from Guangzhou Qian'an Chemical Co., Ltd., China. 1,5,7-Triazabicyclo-[4.4.0]dec-5-ene (TBD), polyethyleneimine (PEI, weight average molecular weight (Mw)=10000 g·mol−1), (3-glycidyloxypropyl) trimethoxysilane (KH560) and zinc acetate (ZA) were purchased from Shanghai Aladdin Biochemical Technology, China. Polyacrylic acid (PAA, Mw=5000 g·mol−1) was obtained from Shanghai Macklin Biochemical Technology Co., Ltd, Shanghai, China. Glass fiber cloth 1080 (electronic grade) was provided by Honghe Electronic Materials Technology Co., Ltd., Shanghai, China. Hexagonal boron nitride (h-BN, lateral size: 2−5 μm, thickness: 200~300 nm) was acquired from Dandong Rijin Technology Co., Ltd., Dandong, China. Boron nitride nanosheets (BNNSs, lateral size: 50 nm) are supplied by Beijing Deke Daojin Science and Technology Co., Ltd., Beijing, China. All the other reagents were purchased from Guangzhou Chemical Reagent Factory and used as received.CharacterizationFourier transform infrared (FTIR) spectra were recorded on a Bruker Tensor 27 from 400 cm−1 to 4000 cm−1. Proton nuclear magnetic resonance (1H-NMR) spectra were obtained via an AVANCE III 400 MHz using chloroform-d as solvent. Chemical elements of the h-BN macron sheet before and after modification was determined using an ESCALab250 X-ray photoelectron spectroscopy. Differential scanning calorimetry (DSC) measurements were conducted on a TA DSC-Q10 in nitrogen atmosphere at a heating rate of 10 °C·min−1. Morphologies and microstructures of the composite laminates were observed by a Hitachi S-4800 scanning electron microscopy (SEM) equipped with an energy dispersive spectrometer (EDS) detector. Tensile tests were detected with an CMT6103 electronic universal testing machine according to ASTM D882 (250 mm × 20 mm × 1 mm) at a cross-head speed of 2 mm·min−1. Electrical breakdown strength was measured following the IPC-TM-650 standard using a CS9915AX Programmable Withstanding Voltage Tester.Dynamic mechanical analysis (DMA) was conducted by a 01 db Metravib DMA25 from 50 °C to 250 °C at a heating rate of 5 °C·min−1. Stress relaxation experiments were performed at a strain of 1% that was applied at t=0. Relaxation activation energy, Ea, was estimated from:[31]1$ \mathrm{l}\mathrm{n}{\tau }^{*}\left(T\right)=\mathrm{l}\mathrm{n}{\tau }_{0}+\frac{{E}_{\mathrm{a}}}{RT} $where τ* is the characteristic relaxation time defined as the time required for the relaxation of initial stress to its 37%; τ0 is the characteristic relaxation time at infinite temperature; R and T are the universal gas constant and the absolute temperature, respectively.Thermal gravimetric analysis (TGA) was carried out on a TGQ50 from 25 °C to 800 °C at a heating rate of 10 °C·min−1 in air atmosphere. The heat-resistance index (THRI) was calculated from:[32]2$ {T}_{\mathrm{H}\mathrm{R}\mathrm{I}}=0.49\times \left[{T}_{5{\text{%}}}+0.6\times ({T}_{30{\text{%}}}-{T}_{5{\text{%}}})\right] $where T5% and T30% represent 5% and 30% weight loss temperature, respectively.The BN content was calculated from:3$ \mathrm{B}\mathrm{N}\;\mathrm{c}\mathrm{o}\mathrm{n}\mathrm{t}\mathrm{e}\mathrm{n}\mathrm{t}=\frac{({m}_{1}-{m}_{2})}{{m}_{0}}\times 100{\text{%}} $where m1 is the residual mass at 800 °C, m2 is the mass of GFC calculated from mass per unit area (46.8 g·m−2) and area of tested sample, and m0 is the sample mass.The through-plane (k⊥) and in-plane (k∥) thermal conductivities were estimated from:4$ k=\alpha \times \rho \times c $where α is the thermal diffusivity measured by an LFA 467 laser-flash diffusivity instrument; ρ is the bulk density detected using an electronic density balance; and c is the specific heat capacity determined by a 200 F3 DSC at a heating rate of 10 °C·min−1.To create artificial interlaminar crack in the composite laminate specimen, a Teflon thin sheet (width=10 mm, thickness=20 μm) was sandwiched between impregnated GFCs and then hot-pressed to cure like common GFREPx composite. Then, the Teflon thin sheet was pulled out to produce an interlaminar crack. Healing of the pre-cracked sample was performed by treating at 230 °C for 3 h under a lateral pressure of 10 MPa. The healing efficiency, η, was estimated from:5$ \eta =\frac{{P}_{\mathrm{h}\mathrm{e}\mathrm{a}\mathrm{l}\mathrm{e}\mathrm{d}}}{{P}_{\mathrm{o}\mathrm{r}\mathrm{i}\mathrm{g}\mathrm{i}\mathrm{n}}} $where Porigin and Phealed are the physical properties (including tensile strength, thermal conductivity and electrical breakdown strength) of the initial and healed samples, respectively.Modification of h-BN Micron PlatesSodium hydroxide (40 g) and deionized water (200 mL) were added into a round bottom flask and magnetically stirred until completely dissolved. Then, h-BN powder (20 g) was dispersed in the alkaline solution, and keep reacting at 120 °C for 24 h. After filtration, the filter residue was washed with water and then dried, yielding h-BN@OH powder bearing active hydroxyl groups.Meantime, ethanol (20 mL) was added to a beaker, followed by the adjusting of pH value to 3−5 using acetic acid. Subsequently, KH560 (2 g) was hydrolyzed in the above-obtained acidic solution for 30 min, and then transferred to the h-BN@OH (20 g)/toluene (40 mL) dispersion solution in a round bottom flask. After reacting at 110 °C for 10 h, the KH560 modified h-BN (h-BN@KH560) was obtained via filtrating, washing with ethanol, and vacuum drying.Preparation of BNNSs Modified Glass Fiber ClothThe BNNS@OH sheets were synthesized following the preparation procedure of h-BN@OH, and then immersed in 5 wt% polycationic electrolyte aqueous solution of PEI at room temperature for 24 h. After filtration and drying, BNNS@PEI carrying positive charges was yielded. Likewise, GFC@PAA bearing negative charges was obtained by soaking in 10 wt% polyanion electrolyte aqueous solution of PAA. Afterwards, BNNS@PEI particles (5 g) were ultrasonically dispersed in distilled water (1 L), and then added to a piece of GFC@PAA immersing in water. After magnetically stirring for 5 h to undergo electrostatic self-assembly between BNNS@PEI and GFC@PAA, GFC@BNNSs with effective in-plane thermally conductive pathways was achieved accordingly.Preparation of Dynamic Reversible Crosslinked Epoxy NetworksTypically, CY-179 and ZAH (molar ratio of ZAH to epoxy group is 0.05/1) were added into a round bottom flask, followed by heating up to 140 °C under mechanical stirring. After cooling to 100 °C, a certain amount of MHHPA hardener was added and mixed evenly (molar ratio of epoxy group to anhydride group is 1.8/1). Subsequently, the reaction mixture was placed in a vacuum oven to remove the bubbles, and then cured in a polytetrafluoroethylene (PTFE) mold at programmed temperature (130 °C, 5 h; 150 °C, 2 h; 180 °C, 2 h; 200 °C, 0.5 h).Fabrication of Healable GFREP CompositeThe dip coating solution was prepared by mixing CY179 (100 g), ZAH (9.2 g) and various concentrations of h-BN@KH560 (i.e., 0 wt%, 5 wt%, 10 wt%, 15 wt%, 25 wt%, 30 wt%) at 140 °C, followed by cooling to 100 °C and then blending with MHHPA (74.0 g).Then, the GFC@BNNSs was impregnated with the dip coating solution by scratch coating, followed by overlapping and hot-pressing with a plate vulcanizer under 15 MPa at 130 °C for 5 h, 150 °C for 2 h, 180 °C for 2 h and 200 °C for 0.5 h, respectively. As a result, series of healable GFREP laminates were produced and labeled as GFREPx (x represents the percentage of h-BN@KH560 in the dynamic epoxy-based dip coating solution).RESULTS AND DISCUSSIONPreparation and Characterization of Dynamic Reversible Crosslinked EpoxyThe synergistic effect between ester bonds and adjacent hydroxy groups, as well as an appropriate catalyst plays an important role in the transesterification reactions and the topological reorganization of dynamic crosslinked epoxy networks.[33] To estimate the thermal stabilities of common transesterification catalysts, TGA curves of ZA, ZAH, and TBD (Fig. 2a) were collected in Figs. 2(b) and 2(c). Among them, ZA possesses superior heat resistance (initial decomposition temperature (Ti)=231.3 °C, maximum weight loss rate temperature (Tp)=275.9 °C), however, poor solubility and high melting temperature (Tm=237 °C) severely limit its application in solvent-free epoxy resin system. In the case of ZAH, the weight loss below 110 °C is ascribed to the associated water molecules, while the decomposition temperature in the second weight loss step (Ti=173.6 °C) is significantly higher than that of TBD (Ti=156.7 °C). On the other hand, the increased peak area of carboxyl moieties in the 1H-NMR spectrum of MHHPA/ZAH mixture pre-treated at 140 °C proves that the associated water molecules of ZAH can even trigger the hydrolysis of anhydride hardener to produce carboxyl moieties, and thus promote the curing reaction of epoxy groups (Fig. 1a). Based on the above-mentioned considerations, ZAH is selected as the catalyst for the subsequent preparation of dynamic reversible crosslinked epoxy and composites with high Tg.Fig 2(a) Chemical structures, (b) TGA curves and (c) DTG curves of various transesterification catalysts; (d) 1H-NMR spectra of MHHPA, ZAH and the MHHPA/ZAH mixture pre-treated at 140 °C for 10 min; (e) Typical FTIR spectrum of the dynamic epoxy network obtained from CY-179 and MHHPA; (f) Normalized stress relaxation curves of dynamic epoxy network measured at various temperatures. The inset shows the fitting of the characteristic relaxation times of dynamic epoxy network at various temperatures according to Arrhenius equation. (g) DSC curve of dynamic epoxy network.As shown in the typical FTIR spectrum of dynamic epoxy network obtained from CY-179 and MHHPA (the molar ratio of epoxy group to anhydride moiety=1.8:1, Fig. 2e), characteristic band of epoxy group (910 cm−1) disappears, proving the completion of curing process. Meantime, the absorption peaks belonging to free hydroxyl group (3352 cm−1) and carboxyl moiety (1733 cm−1) are clearly observed. It indicates that the β-hydroxy ester bonds are generated in the dynamic epoxy networks, offering the structural basis for the intrinsic self-healing capability.To examine the dynamic reversibility of as-fabricated epoxy, stress relaxation experiments were carried out. The normalized stress relaxation curves at different temperatures of dynamic crosslinked epoxy are shown in Fig. 2(f). The relaxation rate enhances with the increase of temperature, owing to the accelerated transesterification rate between included β-hydroxyl ester bonds. Since the relaxation is mainly controlled by the foresaid reversible exchange reactions,[34,35] temperature dependence of the characteristic relaxation time, τ*, can be described by the Arrhenius law. Accordingly, the relaxation activation energy was estimated to be 235.1 kJ·mol−1 by the Eq. (1) (Fig. 2f), which is significantly higher than the values of reported dynamic epoxies containing β-hydroxyl ester bonds (70−120 kJ·mol−1),[36−39] on account of the restricted molecular segment motions induced via the higher Tg (177 °C, Fig. 2g).[37]Modification of h-BN Micron PlatesThe lack of polar moieties on the surface of h-BN micron plates hinder their uniformly dispersion and the formation of continuous heat conduction pathways in the polymers.[40−42] Accordingly, h-BN micron plates were pre-activated by NaOH aqueous solution (h-BN@OH), followed by the chemical grafting of KH560 (h-BN@KH560), in hope of improving the compatibility between h-BN sheets and polymeric matrix. Moreover, epoxy group of KH560 can participate in the curing reactions of the epoxy matrix to enhance the mechanical strength of resultant composites.As shown in the FTIR spectra of h-BN@OH (Fig. 3a), characteristic absorption of hydroxyl group appears at 3400 cm−1 after the chemical treatment of NaOH aqueous solution. When the h-BN@OH was reacted with silane coupling agent (i.e., KH560) via hydrolysis and condensation reactions, the characteristic peak of the hydroxyl moiety disappears. The as-obtained h-BN@KH560 exhibits absorption bands at around 1381 and 816 cm−1, which are attributed to B―N stretching and B―N―B bending, respectively. Meantime, absorption bands belonging to stretching vibrations of C―H (2900 and 2840 cm−1) and Si―O (1100 cm−1) of KH560 are observed.[41] Besides, XPS spectrum of h-BN@KH560 shows a new characteristic peak of Si element at the binding energy of 103 eV (Figs. 3b and 3c).[42] The above results indicate that KH560 was successfully grafted to the surface of h-BN sheets, and the grafting degree is revealed to be 4.6 wt% by TGA (Fig. 3d). In comparation to the morphologies of original h-BN (Fig. 3e), edge boundaries of the h-BN@KH560 micron plates become blurred (Fig. 3f), owing to the grafting of a layer of KH560 modifiers. Even so, no obvious aggregation occurs (Fig. 3g).Fig 3(a) FTIR spectra and (b, c) XPS spectra of h-BN and h-BN@KH560; (d) TGA curve of h-BN and h-BN@KH560; (e−g) SEM images of h-BN (e) before and (f, g) after modification by KH560.Construction of In-plane and Through-plane Thermally Conductive Pathways within the CompositeElectrostatic self-assembly is a facile approach to fabricate hybrid materials for achieving the synergistic improvement effects.[43,44] To construct in-plane thermally conductive pathways and reduce the interfacial thermal resistance, BNNSs were adsorbed onto the surface of GFC via electrostatic self-assembly between PEI coated BNNSs and PAA coated GFC (Fig. 1c). As examined by the FTIR analyses, the PEI and PAA are successfully incorporated to the surfaces of BNNSs (Fig. 4a) and GFC (Fig. 4b), allowing the formation of GFC@BNNSs. The surface topographies of GFC and GFC@BNNSs were characterized by the SEM observations (Figs. 4c−4f). Fig. 4(c) exhibits that the surface of initial GFC is smooth and free of impurities. After electrostatic absorption of BNNSs (Figs. 4d−4f), dense BNNSs particles are attached to the surface of the GFC@BNNSs to form in-plane heat conduction pathways. The loading content of BNNSs is calculated to be only 5.0 wt% by gravimetric method, which avoids the mechanical performance degradation of the composites caused by excessive filler content.Fig 4FTIR spectra of (a) PEI coated BNNSs, and (b) PAA coated GFC. SEM images of the surface of (c) original GFC and (d–f) GFC@BNNSs.Afterwards, the GFC@BNNSs is impregnated with the mixture of h-BN@KH560 micron platelets, CY179 and MHHPA (see the EXPERIMENTAL section for more details) by scratch coating, and then overlapped and hot-pressed to undergo epoxy-anhydride crosslinking reaction at elevated temperature. The BNNSs adsorbed on the GFC surface should be able to closely contact with the h-BN@KH560 platelets dispersed in the epoxy matrix at the assistance of scraper and plate vulcanizer during fabrication procedure. Accordingly, the in-plane and through-plane heat conduction paths are efficiently constructed as expected.To clarify the distribution of BNNSs and h-BN@KH560 micron plates, and the formed heat transport paths in the resultant composite laminate, the morphologies of the cross-sections of the representative GFREP25 are investigated by SEM and element mappings (Fig. 5). No obvious voids and delamination are observed in the region between two layers of GFC@BNNSs (Figs. 5a−5c). Meantime, the BNNSs are densely distributed in the location of GFC and its near surface zones as revealed by the mappings of Si, O and B elements (Figs. 5e−5f). It indicates that the BNNSs electrostatically absorbed on the GFC are closely connected with h-BN@KH560 dispersed in the epoxy matrix without obvious boundaries, which constitutes the through-plane heat conduction pathways.Fig 5(a−c) SEM images, (d) Si, (e) O and (f) B element mappings of the cross-section of GFREP25.Thermal stabilities and BN contents of the healable GFREPx based on dynamic epoxy networks were examined by using TGA (Fig. 6a). The 5% weight loss temperatures (T5%) and heat-resistance indexes (THRI, refer to Eq. 2) of GFREPx composites in air are higher than 360.0 and 200 °C, respectively, showing superior thermal resistance. With the increase of the additive amount of h-BN@KH560 from 5 to 30 wt% in the dynamic epoxy matrix, the total BN content (including BNNSs and BN@KH560) in composites is calculated to be 7.5 wt% to 18.9 wt% using Eq. (3). Accordingly, the Tg value of composite materials significantly enhances from 179 °C to 222 °C measured by DMA (Fig. 6b), on account of the restriction of macromolecular chains mobility.[45,46] Obviously, Tg values of the present composites are higher than that of commercially high-temperature resistant copper clad laminates (about 180 °C). The results imply that the incorporation of reversible bonds does not significantly decay the thermal resistance of the composites.Fig 6(a) TGA curves of GFREPx; (b) tanδ of GFREPx as a function of temperature; (c) Effect of h-BN@KH560 content on thermal conductivities of the composites; (d) Dependences of thermal conductivities of glass fiber cloth reinforced polymer composites on thermally conductive fillers content reported in literature in comparison with the results of this work. Overall thermal conductivity (ko). (e) Tensile properties and (f) volume resistances and breakdown strengths of GFREPx.Afterwards, the thermal conductivities (k∥ and k⊥) of the as-obtained composite materials with different BN contents were analysed by laser-flash method, and the results are shown in Fig. 6(c). GFREP0 excluding h-BN@KH560 sheets in the dynamic epoxy matrix displays k∥ and k⊥ values of 1.36 and 0.27 W·m−1·K−1, respectively. When the loading amount of h-BN@KH560 increases from 5 wt% to 30 wt%, the heat conduction networks within composites become more and more perfect, resulting in the gradual enhancement of thermal conductivities. With respect to the GFREP30, 3.54 and 1.20 W·m−1·K−1 are achieved for the k∥ and k⊥, respectively, which are about 2.6 and 4.4 times those of the GFREP0. The thermally conductive performances of resultant composites are high than most reported values of the glass fiber cloth reinforced polymer composites (Fig. 6d),[6,8−11,47,48] and that of the commercially thermal conductive GFREP (1−1.5 W·m−1·K−1), presenting potential application prospect in the high-end copper clad laminates. Even through the composite Ref. [5] possesses higher thermal conductivities, its mechanical property (tensile strength=80.7 MPa) and insulating characteristic (electrical breakdown strength=18.4 kV·mm−1) have to be sacrificed at high concentration of filler content. Meantime, the lower Tg (153.2 °C) and the built-in Diels-Alder reversible bonds are unfavorable to its application in high-temperature organic substrate.As described above, good mechanical strength and electrical insulation properties are essential for organic substrates during the assembly and usage of PCBs. The tensile strength and elongation at break of the GFREPx slightly deteriorate with the increase of BN content (Fig. 6e), owing to the raise of interface defects between filler and matrix. Meantime, the addition of h-BN fillers also gradually decreases the volume resistance and breakdown strength (Fig. 6f), on account of the increase of leakage current and space charge carriers caused by the interface defects. Nevertheless, tensile strength and elongation at break of 128.7 MPa and 5.0% are determined for the GFREP25 composite at 17.6 wt% BN content and 18.8 wt% GFC. Meantime, high volume resistance and breakdown strength of about 9×1015 Ω·cm and 23.2 kV·mm–1 are remained, which are comparable to the reported data.[43]Healing Capabilities of the CompositeTo investigate the healing capability of composite materials, an interlaminar failure is created as follow. A Teflon film (thickness=20 μm) is embedded between two impregnated GFCs and then hot-pressed to cure, followed by pulling out of the implant (Fig. 7a). Figs. 7(b)−7(e) show that a crack with a width of about 10 μm is made between two pieces of GFCs. Afterwards, the damaged specimen is subjected to healed under 10 MPa at 230 °C for 3 h. Like other reversible crosslinked epoxy networks,[32,49−53] it is found that the crack is completely closed as a result of the dynamic transesterification reactions between fracture surfaces (Figs. 7f−7j). Meantime, the physical recontact of fillers across the crack surface (Fig. 7g) is allowed to restore the continuous heat conduction pathways,[6,21,27] which is confirmed by the uniform distribution of B element around the crack region after healing (Fig. 7j).Fig 7(a) Fabrication of the GFREP specimen with artificial interlaminar damage. SEM images (b, f, g), Si (c, h), O (d, i) and B (e, j) element mappings of the (b−e) damaged and (f−j) healed GFREP25.Moreover, quantitative characterization of healing efficiency is estimated by the restoration of mechanical, electrical insulation, and thermally conductive properties. The interlaminar cracks would not only become the main fracture sources under external force, but also reduce the resistance to penetration by increase the leakage current and space charge carriers. Meantime, the through-plane heat conduction pathways are blocked after the formation of interlaminar damage. As a result, the tensile strength, breakdown strength and k⊥ value are significantly decayed (Fig. 8). Benefitted from the exchange reaction of included β-hydroxyl ester bonds and the accompanied physical recontact of BN sheets, 71.2% of tensile strength, 83.6% of breakdown strength and 69.1% of k⊥ are recovered accordingly (Fig. 8).Fig 8Quantitative characterization of healing efficiencies of CFREP25 calculated from recoveries of tensile strength, electrical breakdown strength and k⊥, respectively.CONCLUSIONSTo develop healable thermally conductive GFREP composite with excellent thermal stability, dynamic crosslinked cycloaliphatic epoxy containing reversible β-hydroxyl ester bonds was compounded with: (i) glass fiber cloth attached by BN nanoparticles via electrostatic self-assembly, and then (ii) micrometer-sized BN platelets through dip coating method, to construct in-plane and through-plane thermal conduction paths, respectively. Accordingly, the as-fabricated GFREP25 composite achieved k∥ and k⊥ of 3.29 and 1.16 W·m−1·K−1 at a total BN content of only 17.6 wt%, which are about 4 and 5 times those of the GFREP composite excluding BN sheets, respectively. Meantime, excellent heat resistance (Tg=204 °C), good mechanical strength (128.7 MPa) and electrical insulating property (23.2 kV·mm−1) were acquired. Besides, decent healing efficiencies of 71.2% (characterized by tensile strength recovery), 83.6% (characterized by breakdown strength) and 69.1% (characterized by thermal conductivity) were measured, respectively. The conflicting properties required for high performance PCBs application, e.g., intrinsic self-healablity, thermal conductivity, mechanical robustness, heat tolerance, were thus successfully united in the current composite material by the proposed strategy.