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Synergistically Improving the Thermal Conductivity and Mechanical Strength of PEEK/MWCNT Nanocomposites by Functionalizing the Matrix with Fluorene Groups

Synergistically Improving the Thermal Conductivity and Mechanical Strength of PEEK/MWCNT... IntroductionEfficient thermal management materials play a crucial role in industrial production and device assembly applications.[1–4] To reduce equipment's thermal aging and malfunction, many critical components are required to have high thermal conductivity to quickly dissipate the heat accumulated during operation. Traditionally, metals are most commonly used thermally conductive materials. Due to the characteristics of low density, high strength, easy processing, high‐temperature resistance, and irradiation resistance, special engineering plastics are gradually replacing traditional metal materials in many fields such as automobile,[5] aerospace,[6] and human bone replacement.[7–9] However, plastics usually possess low intrinsic thermal conductivity (0.25 W m−1 K−1), which severely limits their practical applications in some specialized fields.[10–12] Therefore, the design of filled composites containing carbon or metals has become a significant research interest in the field of thermally conductive materials.As a representative engineering plastic, poly (ether ether ketone) (PEEK) has attracted more and more attention in the research of thermal conductivity modification with its development of high‐temperature applications. Adding thermal conductive fillers to PEEK is undoubtedly one of the most facile and commonly used methods. Lisa et al. prepared PEEK/argentum composites through melt blending, and the addition of 49% silver increased the thermal conductivity of the composite to 0.42 W m−1 K−1.[13] Liu et al. reported boron nitride (BN)/PEEK composites prepared by melt mixing, and the thermal conductivity of the composite reached 1.01 W m−1 K−1 with a loading of 30 wt% BN.[14] However, it can be concluded these micron fillers were usually inefficient that highly filled composites were required to achieve high thermal conductivity, thus inevitably leading to the reduction in other properties of composites, such as mechanical properties.[15–19] Compared with micron fillers, nanofillers (e.g., multi‐walled carbon nanotubes (MWCNT) or graphene) could significantly improve the thermal conductivity of composites with only a small number of additions, while striking a balance between other properties. Hwang et al. prepared PEEK nanocomposites containing graphene oxide (GO)/MWCNT hybrid fillers, and the thermal conductivity can be increased to 0.45 W m−1 K−1 with only 1.5 wt% additions.[20] Li et al. reported the preparation of PEEK/carbon nanotubes (CNT) nanocomposites through melt blending, and the thermal conductivity reached 391% improvement.[21] Although large numbers of studies have demonstrated that the thermal conductivity of PEEK nanocomposites can be improved by adding MWCNT, the reinforcing effect was still far from the results calculated by the theoretical model. This large difference between the experimental results and the theoretical model mainly came from the poor thermal conductivity at the matrix–filler interface, which severely weakened the reinforcing efficiency of thermal conductive nanofillers.[22–25]To solve the problem of poor interface thermal conductivity, considerable work has been done by researchers in improving compatibility and interfacial adhesion.[26–30] Cao et al. prepared poly(vinylidene fluoride)/CNT thermally conductive composites and the thermal conductivity can be improved by 600% through covalent modification of CNT.[27] Guo et al. used triethoxyvinylsilane to chemically modify MWCNT and then prepared poly(vinylidene fluoride)/MWCNT thermally conductive composites. By solving the interface compatibility and agglomeration problems, the interfacial thermal resistance of the composite was reduced by 46%.[31] However, compared with covalent modification, non‐covalent modification of CNT is attracting more attention as it can maintain the high intrinsic thermal conductivity of thermally conductive fillers.[32–34]In this paper, we synthesized a series of PEEK nanocomposites with high MWCNT content via in situ polymerization. A small amount of fluorene conjugated group was introduced into the PEEK chain to strengthen the interfacial non‐covalent bonding, thus reducing the interface thermal resistance between the matrix and fillers. The effects of conjugated fluorene groups on the interfacial thermal resistance and the effects of MWCNT content on the thermal conductivity and mechanical properties of the composites were studied in detail. Considering that the composites can also be used as matrix resins, a multi‐component composite with high thermal conductivity was prepared based on PEEK/MWCNT nanocomposites. This work provides a novel insight into the design of thermally conductive matrix resins and the preparation of multi‐component composites with high thermal conductivity.Experimental SectionMaterials4,4'‐Difluorobenzophenone, hydroquinone, bisphenol fluorene (FD), anhydrous Na2CO3, anhydrous K2CO3, xylene, and diphenyl sulfone (DPS) were obtained from Aladdin (Shanghai, China) and Sinopharm Chemical Reagent (Shanghai, China). The multi‐walled carbon nanotubes (MWCNT) were purchased from Chengdu Organic Chemicals Co., China (length: 5–30 µm, diameter: 10–20 nm). The graphite nanosheet (GNS) was purchased from XF Nano, Inc., China (thickness: less than 40 nm, diameter: 3–6 µm). All monomers and fillers were used without further purification.Synthesis of i‐FD‐PEEK/MWCNT NanocompositesI‐FD‐PEEK/MWCNT nanocomposites were synthesized through tip‐ultrasonic‐assisted in situ copolymerization, as shown in Scheme 1. In detail, diphenyl sulfone was first added into a 1000 mL ringent three‐necked flask and heated to melt at 140 °C, followed by the addition of MWCNT. With a tip ultrasonic device, MWCNT was well untangled and dispersed in solvent DPS by 30 min ultrasonication. 4,4'‐difluorobenzophenone, hydroquinone, bisphenol fluorene, anhydrous K2CO3, and anhydrous Na2CO3 were then added into the suspension. The system was heated to reflux at 170 °C for 2 h to remove the water. After distilling the xylene, the reaction mixture was programmed to 300 °C to polymerize to a black mixture. The mixture was poured into deionized water and washed with acetone and boiling water. After drying at 100 °C for 12 h, the in situ polymerized i‐FDy‐PEEK/xMWCNT (abbreviate to i‐FDy‐PxM) was obtained, where x and y presented the weight amount of MWCNT and molar amount of bisphenol fluorene.1SchemeThe synthesis of i‐PEEK/MWCNT and i‐FD‐PEEK/MWCNT.For i‐PEEK/xMWCNT (abbreviate to i‐PxM), x presented the weight amount of MWCNT and all the bisphenol monomer was hydroquinone. The subsequent reaction steps were the same as i‐FDy‐PEEK/xMWCNT. The reaction flow diagram was shown as Figure 1. Meanwhile, copolymers without MWCNT were synthesized under the same polymerization procedure as references, named FDx‐PEEK.1FigureSchematic representation of in situ copolymerization of i‐PEEK/MWCNT, i‐FD‐PEEK/MWCNT and the multi‐scale segregated network composite.Referring to the previous work,[35] the multi‐scale segregated network composite i‐PEEK/MWCNT@20GNS was prepared by ball‐mill and hot‐press processing. The PEEK/MWCNT core was replaced by i‐FDy‐PEEK/xMWCNT and the composite was named i‐FDy‐PEEK/xMWCNT@20GNS (abbreviate to i‐FDy‐PxM@20GNS).CharacterizationThermal properties and specific heat capacity of nanocomposites were measured by a Mettler Toledo DSC821e instrument from 100 to 400 °C at 10 °C min−1 under nitrogen. The crystallinity was measured by wide‐angle X‐ray diffraction (XRD) using a SHIMADZU LabX XRD‐6000 instrument from 10° to 80° at the rate of 10° min−1. Raman spectrum of MWCNT was performed on a Horiba LabRAM HR Evolution instrument in the range of 1000–2000 cm−1. The chemical structure was characterized by Fourier transform infrared (FT‐IR) spectrometer. Thermal diffusion was carried out with a laser flash apparatus (Netzsch LFA467). The density of the composites was measured by the Archimedes drainage method. The thermal conductivity (λ) was calculated with the formula (1):1λ=α×Cp×ρ\[\begin{array}{*{20}{c}}{\lambda = \alpha \times Cp \times \rho }\end{array}\]where α is the thermal diffusivity (mm2 s−1), Cp is the specific heat capacity (J g−1 K−1), and ρ is the density (g cm−3) of the composite. The mechanical properties of the composites were investigated with a universal testing machine (Shimadzu, AG‐I 1 kN). Storage module (E') was characterized by dynamic mechanical analysis (DMA, TA RSAG2) from 50 to 250 °C at the heating rate of 5 °C min−1. The microtopography was tested by scanning electron microscope (SEM) using a Nova nano 450 instrument. The transmission electron microscopy (TEM) was carried out with a JEM‐2100TEM at 200 kV. The thermogravimetric analysis (TGA) was measured by a METTLER TGA2 instrument from 100 to 800 °C at 10 °C min−1 under air. The solid‐state NMR tests were carried out with a JNM‐ECZ600R spectrometer (JEOL RESONANCE Inc., Japan) (600 MHz). The 13C‐NMR spectra of PEEK and FD‐PEEK were shown in Figure 2b, where 159.97, 151.38, 132.62 121.65, 118.14 ppm were assigned to PEEK and 64.30 ppm was assigned to the quaternary carbon of bisphenol fluorene.2Figurea) Raman spectra of MWCNT before and after ultrasonically dispersed with fluorene and monomers; b) 13C‐NMR curves of pristine PEEK and in situ copolymerization polymers; c) DSC curves and d) XRD pattern of i‐FDy‐P6M with different fluorene additions.Result and DiscussionFirstly, the interaction force between fluorene group and MWCNT was confirmed by Raman spectroscopy test. Figure 2a shows the Raman spectrum of pristine MWCNT, MWCNT ultrasonically dispersed with hydroquinone and 4,4′‐difluoro‐diphenylmethanone (MO@MWCNT), and MWCNT ultrasonically dispersed with bisphenol fluorene (FD@MWCNT). The characteristic peaks of MWCNT at 1349 cm−1 (D peak) and 1577 cm−1 (G peak) were significantly red‐shifted after being coated with bisphenol fluorene, while the peaks of MO@MWCNT were slightly shifted. This indicated stronger π–π interaction between the fluorene‐containing backbone and the carbon nanotubes than that between the fluorene‐free PEEK backbone and the MWCNT, which was more conducive to reducing the interfacial thermal resistance between resin and MWCNT. After ultrasonic dispersion treatment with fluorene, the ID/IG values calculated by the ratio of D peak to G peak remained almost unchanged, indicating that the surface lattice structure of MWCNT was not markedly damaged by ultrasonication.[36] The solid‐state NMR curves of pristine PEEK and i‐FDy‐PxM were shown in Figure 2b. The signals at 159.97, 151.38, 132.62, 121.65, 118.14 ppm were assigned to benzene rings, and quaternary carbon of fluorene was observed at 64.33 ppm, which proved the successful synthesis of fluorene‐containing copolymers. As expected, the peak area of the fluorene group increased with the addition of bisphenol fluorene, calculated by taking the carbonyl peak as the normalization, and the actual copolymerization content was about 4.7 mol% when 5 mol% bisphenol fluorene was added, and about 8.8 mol% when 10 mol% was added. The results indicate that the actual fluorene content generally corresponds to the monomer input. FT‐IR spectra of the fluorene‐containing copolymers were shown in Figure S1 (Supporting Information). It can be observed that the composites had the same fingerprint peaks as those reported in the literature for pristine PEEK.[37]The thermal properties of the nanocomposites with varying fluorene content and a constant amount of MWCNT (6 wt%) were measured by DSC. As shown in Figure 2c, with 2, 5, and 10 mol% fluorene copolymerized, the glass transition temperature (Tg) of i‐FDy‐P6M nanocomposites increased from 145 to 149, 153, and 162 °C, respectively. As a contrast, Tg of FDx‐PEEK copolymers without MWCNT only increased from 143 to 145,148, and 152 °C, respectively, as shown in Figure S2 and Table S1 (Supporting Information). The difference in the Tg increase was another proof that the fluorene groups with conjugated structure would generate π–π interaction with MWCNT, which would further limit the movement of the copolymer chains, thereby increasing the glass transition temperature.[34] At the same time, it could be seen from the DSC curves of the nanocomposites that the cold crystallization peak temperatures (Tc) also increased significantly with the increasing fluorene content. This may be because the stronger interaction between the higher fluorene‐containing chain and MWCNT would further reduce the regular arrangement ability of the polymer chain, thus a higher temperature was required to promote the regular arrangement occurring. It should also be noted that both the melting point temperatures (Tm) and melting enthalpy of the nanocomposites slightly decreased with the increase of fluorene content. This was attributed to the chain regularity disrupted by the addition of fluorene groups, which affected the crystallization properties of the polymer. The crystalline properties of the nanocomposites were characterized by XRD and the spectra were shown in Figure 2d. It can be observed that the characteristic diffraction peaks of the nanocomposites located at 18.9°(110), 20.3°(111), 23.4°(200), 29.0°(211), and 26.1°(002).[38] It demonstrated that the crystal form of the copolymers was not affected by the introduction of fluorene groups, and still displayed the characteristics of semi‐crystalline polymers. At the same time, the fitting calculation showed in Table S1 (Supporting Information) that the crystallinity of the copolymers decreased significantly with the increasing fluorene content from 34.3% to 16.1%, which was consistent with the DSC calculation results.To characterize the interfacial bonding between MWCNT and the matrix, we performed cryosections and TEM tests on the nanocomposites. It can be observed from Figure 3a that the surface of MWCNT was smooth and the boundaries were clear, indicating the poor interaction force between MWCNT and the polymer matrix. With 2 mol% fluorene groups added, the interfacial bonding force was improved through π–π interaction.[34] Therefore, the copolymers were coated on the surface of MWCNT, as observed in Figure 3b. By comparing the SEM images of the quenched sections of composites shown in Figure 3c,d and the enlarged images in Figure 3e,f, it was observed that the surface of MWCNT in i‐P10M was smooth, with almost no polymer cladding visible. This indicates a poor interface between the polymer and MWCNT, leading to inadequate dispersion of MWCNT in the fluorene‐absent polymer matrix. In contrast, the MWCNT in i‐FD2‐P10M was found to be covered by the resin layer, indicating a good interfacial bond between them. Additionally, TEM tests were conducted on i‐FD2‐P10M samples, as shown in Figure S3 (Supporting Information), which demonstrated a uniform dispersion of a large amount of MWCNT in the matrix resin.3Figurea,c) TEM and SEM images of i‐PEEK/10MWCNT, b,d) i‐FD2‐PEEK/10MWCNT, and e,f) the enlarged images.The thermal conductivity of the nanocomposites with varying fluorene content and a constant amount of MWCNT (6 wt%) was tested and shown in Figure 4a. It could be seen that the thermal conductivity of the nanocomposites significantly improved with a small amount of fluorene introduction. With 2 mol% fluorene introduced, the thermal conductivity of i‐FD2‐P6M reached 1.52 W m−1 K−1, which was 27% higher than that without fluorene. This can be concluded that the introduction of a small amount of fluorene improved the interfacial adhesion between the matrix and MWCNT, thereby improving the efficiency of heat transfer. However, the continued increase of the fluorene content did not help to remarkably improve the thermal conductivity and even decreased it when reaching 10 mol%. This can be attributed to the reduced intrinsic thermal conductivity of the copolymers by the excessive introduction of fluorene groups, as shown in Figure 4b. Excessive fluorene groups disrupted the regularity of the molecular structure and weakening the crystallization ability of the polymer chain, thereby reducing the intrinsic thermal conductivity. At present, many studies have found that the interface thermal resistance can be calculated by a specific model, so we also calculated the interface thermal resistance between the copolymers and MWCNT by the formula shown in Equations (2––5).[39,40]2Ke= Km3+f(βx+βz)2−fβx\[\begin{array}{*{20}{c}}{{K_{\rm{e}}} = \;{K_{\rm{m}}}\frac{{3 + f\left( {{\beta _{\rm{x}}} + {\beta _{\rm{z}}}} \right)}}{{2 - f{\beta _{\rm{x}}}}}}\end{array}\]3βx =2(K11c−Km)K11c+Km, βz =K33cKm −1\[\begin{array}{*{20}{c}}{{\beta _{\rm{x}}}\; = \frac{{2\left( {K_{11}^{\rm{c}} - {K_{\rm{m}}}} \right)}}{{K_{11}^{\rm{c}} + {K_{\rm{m}}}}},\:{\beta _{\rm{z}}}\; = \frac{{K_{33}^{\rm{c}}}}{{{K_{\rm{m}}}}}\; - 1}\end{array}\]4K11c=Kc1+2aKdKcKm,  K33C=Kc1+2aKLKcKm\[\begin{array}{*{20}{c}}{K_{11}^{\rm{c}} = \frac{{{K_{\rm{c}}}}}{{1 + \frac{{2{a_K}}}{d}\frac{{{K_{\rm{c}}}}}{{{K_{\rm{m}}}}}}},\;\:K_{33}^{\rm{C}} = \frac{{{K_{\rm{c}}}}}{{1 + \frac{{2{a_K}}}{L}\frac{{{K_{\rm{c}}}}}{{{K_{\rm{m}}}}}}}}\end{array}\]5aK=RK Km\[\begin{array}{*{20}{c}}{{a_{\rm{K}}} = {R_{\rm{K}}}\;{K_{\rm{m}}}}\end{array}\]Ke, Kc, and Km are the thermal conductivity of the nanocomposite, MWCNT and the polymer matrix. The value of 3000 W m−1 K−1 is used for Kc[22] L is the length of MWCNT, d is the diameter of MWCNT, and f is the volume fraction, while in this case d is 10 nm and L is 20 µm. aK is the Kapitza radius. K11C$K_{11}^{\rm{C}}$ and K33C$K_{33}^{\rm{C}}$ are, respectively, the equivalent thermal conductivities along transverse and longitudinal axes of a composite unit cell, i.e., a nanotube coated with a very thin interfacial thermal barrier. βX and βZ are the thermal conductivity enhancement effects of carbon nanotubes compared to the matrix material in the x and z directions, respectively. The Km value is changing with different fluorene addition. By calculation, the interface thermal resistance of pristine PEEK and MWCNT was 1.2E‐8 m2 K W−1, while with 2 and 5 mol% FD fluorene content added, the interface thermal resistance of i‐FD2‐P6M and i‐FD5‐P6M was reduced to 1.0E‐8 and 0.9E‐8 m2 K W−1, respectively. It demonstrated that the introduction of fluorene was effective in reducing the interfacial thermal resistance between the matrix and the filler. When the fluorene content continued to increase to 10 mol%, the thermal conductivity of the nanocomposite started to decrease and the thermal resistance rose to 1.4 E‐8 m2 K W−1. The calculated interface thermal resistance was consistent with the experimental thermal conductivity of the nanocomposites, as shown in Figure 4a,c.4Figurea) Thermal conductivity of i‐FDy‐P6M with different fluorene content; b) thermal conductivity of FDy‐PEEK with different fluorene content; c) interface resistance of i‐FDy‐P6M with different fluorene content; d) tensile strength and elongation at break of i‐FDy‐P6M with different fluorene addition.The mechanical properties of nanocomposites were measured and shown in Figure 4d. The tensile strength of the composites showed a trend of first rising and then falling. When a small amount of fluorene group was introduced, although the crystallinity of the polymer was slightly reduced, the increased bonding force between the resin and the filler would make up for the loss. This is why the tensile strength of i‐FD2‐P6M was enhanced to 106.5 MPa with 2 mol% fluorene groups introduced, which was 10.4% higher than that of i‐P6M. However, continuing to increase the content of fluorene groups would further limit the crystallization ability of the copolymers. The drop in crystallinity led to a significant decrease in the intrinsic strength of the polymer matrix, which cannot even be compensated by the reinforced interfacial bonding. Therefore, the tensile strength of i‐FD5‐P6M started to decrease. The elongation at break of the nanocomposites showed the same trend. With the rising fluorene content, the elongation at break of the nanocomposites increased significantly first and then decreased, reaching the maximum of 8.5% when 2 mol% fluorene was introduced.[34,41–43] A combined consideration of the thermal conductivity and mechanical properties gives that an introduction of 2 mol% fluorene groups is the optimal ratio for our in situ polymerization method to prepare the fluorene‐functionalized PEEK/MWCNT nanocomposites.With the optimal fluorene content determined, the nanocomposites with or without fluorene groups at various MWCNT amounts were selected to be compared. Figure 5a shows the thermal conductivity of i‐PxM without fluorene introduced and i‐FD2‐PxM with 2 mol% fluorene introduced. The thermal conductivity of i‐FD2‐PxM always showed an increment after introducing fluorene and reached 2.41 W m−1 K−1 for i‐FD2‐P10M. Meanwhile, the tensile strength was compared in Figure 5b, showing that the introduction of fluorene allowed the tensile strength of i‐FD2‐P10M to reach 104.3 MPa and the elongation at break increased to 6.4%. In order to elucidate the enhanced the elongation at break of the i‐FD2‐P10M composite, its tensile section was examined using scanning electron microscopy and was presented in Figure 5c,d. As observed from the SEM images, the i‐P10M section exhibited significant agglomeration, leading to stress concentration and a reduction in elongation at break. However, with 2% fluorene group introduced into the polymer chain, the tensile section of the composite became uneven, with the morphology of the resin stretched, and the length of the MWCNT being pulled out of the resin increased significantly. This improvement can be attributed to the significant enhancement in the interfacial bonding between the resin matrix and carbon nanotubes due to the introduction of fluorene groups, leading to better stress transfer to MWCNT during tensile processes and thus improving the strength and elongation at break of the nanocomposites.5Figurea) Thermal conductivity of i‐PxM and i‐FD2‐PxM with different MWCNT additions; b) tensile properties of i‐P10M and i‐FD2‐P10M; c) tensile fracture surface morphology of i‐P10M; and d) i‐FD2‐P10M; e) elevated temperature thermal conductivity of i‐FD2‐P10M; and f) DMA curves of i‐P10 and i‐FD2‐P10M.As high‐temperature resistant materials, the PEEK nanocomposites are usually required to be evaluated at high temperatures. Thermogravimetric analysis was performed on the nanocomposite i‐P10M and i‐FD2‐P10M, and the results are presented in Figure S4 (Supporting Information). It was found that when the fluorene group was incorporated, the composites exhibited thermal stability with a 5% thermal decomposition temperature above 500 °C. This indicates that i‐FD2‐P10M composite can meet the severe processing temperature. The thermal conductivity of i‐FD2‐P10M at elevatored temperatures was measured and shown in Figure 5e. With the increasing temperatures, the thermal conductivity of the nanocomposite gradually rose and reached 2.96 W m−1 K−1 at 150 °C (a temperature close to Tg). Although high temperature would increase the interfacial scattering of phonons, it also improved the intrinsic heat transfer efficiency inside the carbon material.[44] Therefore, the thermal conductivity of the nanocomposite tended to increase with the rising temperature. This showed that the materials still had good performance at high‐temperature conditions. In addition, the DMA measurement was carried out on i‐P10M and i‐FD2‐P10M and their curves were shown in Figure 5f. The prepared nanocomposites maintained high storage modulus (G') before 150 °C. Benefiting from the better interface interaction, G’ of i‐FD2‐P10M was slightly higher than that of i‐P10M. In consistent with DSC test results, i‐FD2‐P10M had a higher glass transition temperature than i‐P10M. To conclude, the introduction of fluorene groups enhanced the interface between the polymer matrix and MWCNT, further limiting the molecular chain movement and improving the heat resistance. After glass transition temperature, G’ of i‐FD2‐P10M was slightly lower than that of i‐P10M, which was mainly affected by the crystallinity of the copolymers.[45]Nanocomposites can also be used as a matrix to prepare multi‐component composites. Following our previous work, a double segregated network composite i‐PEEK/MWCNT@GNS was prepared by replacing the matrix with the fluorene‐functionalized PEEK/MWCNT nanocomposite.[35] As shown in Figure 6, after introducing fluorene groups, the thermal conductivity of i‐FD2‐P10M@20GNS was improved to 3.91 W m−1 K−1, which exhibited a significant increment compared with that of i‐P10M@20GNS. At the same time, the tensile property of the composite was also remarkably improved from 57.7 MPa to 65.8 MPa. The ternary composite's strength enhancement was due to the improved intrinsic strength of the fluorene‐functionalized PEEK/MWCNT nanocomposite matrix. Furthermore, in comparison with other PEEK/fillers thermal conductive composites reported in the literature as shown in Figure 6 and Figure S5 (Supporting Information), the thermal conductivity in our work reached a relatively high level.[13,14,20,46–54] The combination of in situ polymerization and molecular structure design improved the dispersion state of carbon nanotubes, introduced a large amount of MWCNT as a thermally conductive filler, and significantly improved the thermal conductivity interface, resulting in better thermal conductivity of the synthesized nanocomposites. This approach provides a novel view to design multi‐component composites with balanced thermal conductivity and mechanical strength.6Figurea) Comprehensive performance chart of different matrix composite 30 wt% filler; b) comparison of thermal conductivity between the composites in this work and other PEEK/fillers composites reported by literature (this work‐1 represents i‐FD2‐P10M and this work‐2 represents i‐FD2‐P10M@20GNS).To visually demonstrate the heat dissipation capability of the composites, a device was constructed by integrating the composite with an LED light and the heat dissipation effect was tested under the actual working scenario, which was shown in Figure 7a. The real‐time temperature of the composites above the LED was tested by infrared thermography, and the curve was plotted in Figure 7b. The temperature of the material gradually increased with working time, but the surface temperature of i‐FD2‐P10M was consistently lower than the other two references. The final temperature of i‐FD2‐P10M was 17 °C lower than that of pristine PEEK, indicating that i‐FD2‐P10M has superior heat dissipation ability and can significantly reduce the operating temperature of LED devices.7FigureDiagram of a) LED heat sink parts, b) operating temperature curves, and c) infrared thermogram.ConclusionIn the present work, the fluorene‐functionalized PEEK/MWCNT nanocomposites with high thermal conductivity and tensile strength were synthesized through an in situ copolymerization procedure. The introduction of fluorene groups into PEEK can effectively improve the interfacial bonding between the copolymer matrix and MWCNT, thus reducing the interface thermal resistance. The thermal conductivity of the nanocomposite i‐FD2‐P10M with 2 mol% fluorene and 10 wt% MWCNT is increased to 2.41 W m−1 K−1, which is 868% higher than pristine PEEK. Meanwhile, the tensile strength of i‐FD2‐P10M maintain at 104.3 MPa. Based on this nanocomposite matrix, multi‐component composites with segregated network structure are constructed to achieve promising thermal management materials with balanced thermal conductivity and mechanical strength properties. This work provided a generic method to balance the thermal conductivity and mechanical strength of the polymer‐matrix composites and to prepare highly nano‐filled composites with optimized interface adhesion.AcknowledgementsThe authors would like to thank the National and Local Joint Engineering Laboratory for the Synthesis Technology of High‐Performance Polymers for their assistance and use of the facility. This work was financially supported by the Department of Science and Technology of Jilin Province (Grant 20210509034RQ and 20210201117GX) and the Fundamental Research Funds for the Central Universities, the National Natural Science Foundation of China (No. 52003099), the Natural Science Foundation of Jilin Province (YDZJ202101ZYTS002), and the Capital Construction Fund of Jilin Province (2021C039‐1).Conflict of InterestThe authors declare no conflict of interest.Author ContributionsZ.J. and B.J. contributed equally to this work and should be considered co‐first authors. Conceptualization, methodology, investigation, writing original draft. B.J.: Validation, resources, writing original draft. B.Y.: Writing—review and editing, validation. X.L.: Resources, validation. Y.Y.: Investigation, validation. Y.S.: Validation, data curation, writing, reviewing and editing. 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Synergistically Improving the Thermal Conductivity and Mechanical Strength of PEEK/MWCNT Nanocomposites by Functionalizing the Matrix with Fluorene Groups

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2196-7350
DOI
10.1002/admi.202202508
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Abstract

IntroductionEfficient thermal management materials play a crucial role in industrial production and device assembly applications.[1–4] To reduce equipment's thermal aging and malfunction, many critical components are required to have high thermal conductivity to quickly dissipate the heat accumulated during operation. Traditionally, metals are most commonly used thermally conductive materials. Due to the characteristics of low density, high strength, easy processing, high‐temperature resistance, and irradiation resistance, special engineering plastics are gradually replacing traditional metal materials in many fields such as automobile,[5] aerospace,[6] and human bone replacement.[7–9] However, plastics usually possess low intrinsic thermal conductivity (0.25 W m−1 K−1), which severely limits their practical applications in some specialized fields.[10–12] Therefore, the design of filled composites containing carbon or metals has become a significant research interest in the field of thermally conductive materials.As a representative engineering plastic, poly (ether ether ketone) (PEEK) has attracted more and more attention in the research of thermal conductivity modification with its development of high‐temperature applications. Adding thermal conductive fillers to PEEK is undoubtedly one of the most facile and commonly used methods. Lisa et al. prepared PEEK/argentum composites through melt blending, and the addition of 49% silver increased the thermal conductivity of the composite to 0.42 W m−1 K−1.[13] Liu et al. reported boron nitride (BN)/PEEK composites prepared by melt mixing, and the thermal conductivity of the composite reached 1.01 W m−1 K−1 with a loading of 30 wt% BN.[14] However, it can be concluded these micron fillers were usually inefficient that highly filled composites were required to achieve high thermal conductivity, thus inevitably leading to the reduction in other properties of composites, such as mechanical properties.[15–19] Compared with micron fillers, nanofillers (e.g., multi‐walled carbon nanotubes (MWCNT) or graphene) could significantly improve the thermal conductivity of composites with only a small number of additions, while striking a balance between other properties. Hwang et al. prepared PEEK nanocomposites containing graphene oxide (GO)/MWCNT hybrid fillers, and the thermal conductivity can be increased to 0.45 W m−1 K−1 with only 1.5 wt% additions.[20] Li et al. reported the preparation of PEEK/carbon nanotubes (CNT) nanocomposites through melt blending, and the thermal conductivity reached 391% improvement.[21] Although large numbers of studies have demonstrated that the thermal conductivity of PEEK nanocomposites can be improved by adding MWCNT, the reinforcing effect was still far from the results calculated by the theoretical model. This large difference between the experimental results and the theoretical model mainly came from the poor thermal conductivity at the matrix–filler interface, which severely weakened the reinforcing efficiency of thermal conductive nanofillers.[22–25]To solve the problem of poor interface thermal conductivity, considerable work has been done by researchers in improving compatibility and interfacial adhesion.[26–30] Cao et al. prepared poly(vinylidene fluoride)/CNT thermally conductive composites and the thermal conductivity can be improved by 600% through covalent modification of CNT.[27] Guo et al. used triethoxyvinylsilane to chemically modify MWCNT and then prepared poly(vinylidene fluoride)/MWCNT thermally conductive composites. By solving the interface compatibility and agglomeration problems, the interfacial thermal resistance of the composite was reduced by 46%.[31] However, compared with covalent modification, non‐covalent modification of CNT is attracting more attention as it can maintain the high intrinsic thermal conductivity of thermally conductive fillers.[32–34]In this paper, we synthesized a series of PEEK nanocomposites with high MWCNT content via in situ polymerization. A small amount of fluorene conjugated group was introduced into the PEEK chain to strengthen the interfacial non‐covalent bonding, thus reducing the interface thermal resistance between the matrix and fillers. The effects of conjugated fluorene groups on the interfacial thermal resistance and the effects of MWCNT content on the thermal conductivity and mechanical properties of the composites were studied in detail. Considering that the composites can also be used as matrix resins, a multi‐component composite with high thermal conductivity was prepared based on PEEK/MWCNT nanocomposites. This work provides a novel insight into the design of thermally conductive matrix resins and the preparation of multi‐component composites with high thermal conductivity.Experimental SectionMaterials4,4'‐Difluorobenzophenone, hydroquinone, bisphenol fluorene (FD), anhydrous Na2CO3, anhydrous K2CO3, xylene, and diphenyl sulfone (DPS) were obtained from Aladdin (Shanghai, China) and Sinopharm Chemical Reagent (Shanghai, China). The multi‐walled carbon nanotubes (MWCNT) were purchased from Chengdu Organic Chemicals Co., China (length: 5–30 µm, diameter: 10–20 nm). The graphite nanosheet (GNS) was purchased from XF Nano, Inc., China (thickness: less than 40 nm, diameter: 3–6 µm). All monomers and fillers were used without further purification.Synthesis of i‐FD‐PEEK/MWCNT NanocompositesI‐FD‐PEEK/MWCNT nanocomposites were synthesized through tip‐ultrasonic‐assisted in situ copolymerization, as shown in Scheme 1. In detail, diphenyl sulfone was first added into a 1000 mL ringent three‐necked flask and heated to melt at 140 °C, followed by the addition of MWCNT. With a tip ultrasonic device, MWCNT was well untangled and dispersed in solvent DPS by 30 min ultrasonication. 4,4'‐difluorobenzophenone, hydroquinone, bisphenol fluorene, anhydrous K2CO3, and anhydrous Na2CO3 were then added into the suspension. The system was heated to reflux at 170 °C for 2 h to remove the water. After distilling the xylene, the reaction mixture was programmed to 300 °C to polymerize to a black mixture. The mixture was poured into deionized water and washed with acetone and boiling water. After drying at 100 °C for 12 h, the in situ polymerized i‐FDy‐PEEK/xMWCNT (abbreviate to i‐FDy‐PxM) was obtained, where x and y presented the weight amount of MWCNT and molar amount of bisphenol fluorene.1SchemeThe synthesis of i‐PEEK/MWCNT and i‐FD‐PEEK/MWCNT.For i‐PEEK/xMWCNT (abbreviate to i‐PxM), x presented the weight amount of MWCNT and all the bisphenol monomer was hydroquinone. The subsequent reaction steps were the same as i‐FDy‐PEEK/xMWCNT. The reaction flow diagram was shown as Figure 1. Meanwhile, copolymers without MWCNT were synthesized under the same polymerization procedure as references, named FDx‐PEEK.1FigureSchematic representation of in situ copolymerization of i‐PEEK/MWCNT, i‐FD‐PEEK/MWCNT and the multi‐scale segregated network composite.Referring to the previous work,[35] the multi‐scale segregated network composite i‐PEEK/MWCNT@20GNS was prepared by ball‐mill and hot‐press processing. The PEEK/MWCNT core was replaced by i‐FDy‐PEEK/xMWCNT and the composite was named i‐FDy‐PEEK/xMWCNT@20GNS (abbreviate to i‐FDy‐PxM@20GNS).CharacterizationThermal properties and specific heat capacity of nanocomposites were measured by a Mettler Toledo DSC821e instrument from 100 to 400 °C at 10 °C min−1 under nitrogen. The crystallinity was measured by wide‐angle X‐ray diffraction (XRD) using a SHIMADZU LabX XRD‐6000 instrument from 10° to 80° at the rate of 10° min−1. Raman spectrum of MWCNT was performed on a Horiba LabRAM HR Evolution instrument in the range of 1000–2000 cm−1. The chemical structure was characterized by Fourier transform infrared (FT‐IR) spectrometer. Thermal diffusion was carried out with a laser flash apparatus (Netzsch LFA467). The density of the composites was measured by the Archimedes drainage method. The thermal conductivity (λ) was calculated with the formula (1):1λ=α×Cp×ρ\[\begin{array}{*{20}{c}}{\lambda = \alpha \times Cp \times \rho }\end{array}\]where α is the thermal diffusivity (mm2 s−1), Cp is the specific heat capacity (J g−1 K−1), and ρ is the density (g cm−3) of the composite. The mechanical properties of the composites were investigated with a universal testing machine (Shimadzu, AG‐I 1 kN). Storage module (E') was characterized by dynamic mechanical analysis (DMA, TA RSAG2) from 50 to 250 °C at the heating rate of 5 °C min−1. The microtopography was tested by scanning electron microscope (SEM) using a Nova nano 450 instrument. The transmission electron microscopy (TEM) was carried out with a JEM‐2100TEM at 200 kV. The thermogravimetric analysis (TGA) was measured by a METTLER TGA2 instrument from 100 to 800 °C at 10 °C min−1 under air. The solid‐state NMR tests were carried out with a JNM‐ECZ600R spectrometer (JEOL RESONANCE Inc., Japan) (600 MHz). The 13C‐NMR spectra of PEEK and FD‐PEEK were shown in Figure 2b, where 159.97, 151.38, 132.62 121.65, 118.14 ppm were assigned to PEEK and 64.30 ppm was assigned to the quaternary carbon of bisphenol fluorene.2Figurea) Raman spectra of MWCNT before and after ultrasonically dispersed with fluorene and monomers; b) 13C‐NMR curves of pristine PEEK and in situ copolymerization polymers; c) DSC curves and d) XRD pattern of i‐FDy‐P6M with different fluorene additions.Result and DiscussionFirstly, the interaction force between fluorene group and MWCNT was confirmed by Raman spectroscopy test. Figure 2a shows the Raman spectrum of pristine MWCNT, MWCNT ultrasonically dispersed with hydroquinone and 4,4′‐difluoro‐diphenylmethanone (MO@MWCNT), and MWCNT ultrasonically dispersed with bisphenol fluorene (FD@MWCNT). The characteristic peaks of MWCNT at 1349 cm−1 (D peak) and 1577 cm−1 (G peak) were significantly red‐shifted after being coated with bisphenol fluorene, while the peaks of MO@MWCNT were slightly shifted. This indicated stronger π–π interaction between the fluorene‐containing backbone and the carbon nanotubes than that between the fluorene‐free PEEK backbone and the MWCNT, which was more conducive to reducing the interfacial thermal resistance between resin and MWCNT. After ultrasonic dispersion treatment with fluorene, the ID/IG values calculated by the ratio of D peak to G peak remained almost unchanged, indicating that the surface lattice structure of MWCNT was not markedly damaged by ultrasonication.[36] The solid‐state NMR curves of pristine PEEK and i‐FDy‐PxM were shown in Figure 2b. The signals at 159.97, 151.38, 132.62, 121.65, 118.14 ppm were assigned to benzene rings, and quaternary carbon of fluorene was observed at 64.33 ppm, which proved the successful synthesis of fluorene‐containing copolymers. As expected, the peak area of the fluorene group increased with the addition of bisphenol fluorene, calculated by taking the carbonyl peak as the normalization, and the actual copolymerization content was about 4.7 mol% when 5 mol% bisphenol fluorene was added, and about 8.8 mol% when 10 mol% was added. The results indicate that the actual fluorene content generally corresponds to the monomer input. FT‐IR spectra of the fluorene‐containing copolymers were shown in Figure S1 (Supporting Information). It can be observed that the composites had the same fingerprint peaks as those reported in the literature for pristine PEEK.[37]The thermal properties of the nanocomposites with varying fluorene content and a constant amount of MWCNT (6 wt%) were measured by DSC. As shown in Figure 2c, with 2, 5, and 10 mol% fluorene copolymerized, the glass transition temperature (Tg) of i‐FDy‐P6M nanocomposites increased from 145 to 149, 153, and 162 °C, respectively. As a contrast, Tg of FDx‐PEEK copolymers without MWCNT only increased from 143 to 145,148, and 152 °C, respectively, as shown in Figure S2 and Table S1 (Supporting Information). The difference in the Tg increase was another proof that the fluorene groups with conjugated structure would generate π–π interaction with MWCNT, which would further limit the movement of the copolymer chains, thereby increasing the glass transition temperature.[34] At the same time, it could be seen from the DSC curves of the nanocomposites that the cold crystallization peak temperatures (Tc) also increased significantly with the increasing fluorene content. This may be because the stronger interaction between the higher fluorene‐containing chain and MWCNT would further reduce the regular arrangement ability of the polymer chain, thus a higher temperature was required to promote the regular arrangement occurring. It should also be noted that both the melting point temperatures (Tm) and melting enthalpy of the nanocomposites slightly decreased with the increase of fluorene content. This was attributed to the chain regularity disrupted by the addition of fluorene groups, which affected the crystallization properties of the polymer. The crystalline properties of the nanocomposites were characterized by XRD and the spectra were shown in Figure 2d. It can be observed that the characteristic diffraction peaks of the nanocomposites located at 18.9°(110), 20.3°(111), 23.4°(200), 29.0°(211), and 26.1°(002).[38] It demonstrated that the crystal form of the copolymers was not affected by the introduction of fluorene groups, and still displayed the characteristics of semi‐crystalline polymers. At the same time, the fitting calculation showed in Table S1 (Supporting Information) that the crystallinity of the copolymers decreased significantly with the increasing fluorene content from 34.3% to 16.1%, which was consistent with the DSC calculation results.To characterize the interfacial bonding between MWCNT and the matrix, we performed cryosections and TEM tests on the nanocomposites. It can be observed from Figure 3a that the surface of MWCNT was smooth and the boundaries were clear, indicating the poor interaction force between MWCNT and the polymer matrix. With 2 mol% fluorene groups added, the interfacial bonding force was improved through π–π interaction.[34] Therefore, the copolymers were coated on the surface of MWCNT, as observed in Figure 3b. By comparing the SEM images of the quenched sections of composites shown in Figure 3c,d and the enlarged images in Figure 3e,f, it was observed that the surface of MWCNT in i‐P10M was smooth, with almost no polymer cladding visible. This indicates a poor interface between the polymer and MWCNT, leading to inadequate dispersion of MWCNT in the fluorene‐absent polymer matrix. In contrast, the MWCNT in i‐FD2‐P10M was found to be covered by the resin layer, indicating a good interfacial bond between them. Additionally, TEM tests were conducted on i‐FD2‐P10M samples, as shown in Figure S3 (Supporting Information), which demonstrated a uniform dispersion of a large amount of MWCNT in the matrix resin.3Figurea,c) TEM and SEM images of i‐PEEK/10MWCNT, b,d) i‐FD2‐PEEK/10MWCNT, and e,f) the enlarged images.The thermal conductivity of the nanocomposites with varying fluorene content and a constant amount of MWCNT (6 wt%) was tested and shown in Figure 4a. It could be seen that the thermal conductivity of the nanocomposites significantly improved with a small amount of fluorene introduction. With 2 mol% fluorene introduced, the thermal conductivity of i‐FD2‐P6M reached 1.52 W m−1 K−1, which was 27% higher than that without fluorene. This can be concluded that the introduction of a small amount of fluorene improved the interfacial adhesion between the matrix and MWCNT, thereby improving the efficiency of heat transfer. However, the continued increase of the fluorene content did not help to remarkably improve the thermal conductivity and even decreased it when reaching 10 mol%. This can be attributed to the reduced intrinsic thermal conductivity of the copolymers by the excessive introduction of fluorene groups, as shown in Figure 4b. Excessive fluorene groups disrupted the regularity of the molecular structure and weakening the crystallization ability of the polymer chain, thereby reducing the intrinsic thermal conductivity. At present, many studies have found that the interface thermal resistance can be calculated by a specific model, so we also calculated the interface thermal resistance between the copolymers and MWCNT by the formula shown in Equations (2––5).[39,40]2Ke= Km3+f(βx+βz)2−fβx\[\begin{array}{*{20}{c}}{{K_{\rm{e}}} = \;{K_{\rm{m}}}\frac{{3 + f\left( {{\beta _{\rm{x}}} + {\beta _{\rm{z}}}} \right)}}{{2 - f{\beta _{\rm{x}}}}}}\end{array}\]3βx =2(K11c−Km)K11c+Km, βz =K33cKm −1\[\begin{array}{*{20}{c}}{{\beta _{\rm{x}}}\; = \frac{{2\left( {K_{11}^{\rm{c}} - {K_{\rm{m}}}} \right)}}{{K_{11}^{\rm{c}} + {K_{\rm{m}}}}},\:{\beta _{\rm{z}}}\; = \frac{{K_{33}^{\rm{c}}}}{{{K_{\rm{m}}}}}\; - 1}\end{array}\]4K11c=Kc1+2aKdKcKm,  K33C=Kc1+2aKLKcKm\[\begin{array}{*{20}{c}}{K_{11}^{\rm{c}} = \frac{{{K_{\rm{c}}}}}{{1 + \frac{{2{a_K}}}{d}\frac{{{K_{\rm{c}}}}}{{{K_{\rm{m}}}}}}},\;\:K_{33}^{\rm{C}} = \frac{{{K_{\rm{c}}}}}{{1 + \frac{{2{a_K}}}{L}\frac{{{K_{\rm{c}}}}}{{{K_{\rm{m}}}}}}}}\end{array}\]5aK=RK Km\[\begin{array}{*{20}{c}}{{a_{\rm{K}}} = {R_{\rm{K}}}\;{K_{\rm{m}}}}\end{array}\]Ke, Kc, and Km are the thermal conductivity of the nanocomposite, MWCNT and the polymer matrix. The value of 3000 W m−1 K−1 is used for Kc[22] L is the length of MWCNT, d is the diameter of MWCNT, and f is the volume fraction, while in this case d is 10 nm and L is 20 µm. aK is the Kapitza radius. K11C$K_{11}^{\rm{C}}$ and K33C$K_{33}^{\rm{C}}$ are, respectively, the equivalent thermal conductivities along transverse and longitudinal axes of a composite unit cell, i.e., a nanotube coated with a very thin interfacial thermal barrier. βX and βZ are the thermal conductivity enhancement effects of carbon nanotubes compared to the matrix material in the x and z directions, respectively. The Km value is changing with different fluorene addition. By calculation, the interface thermal resistance of pristine PEEK and MWCNT was 1.2E‐8 m2 K W−1, while with 2 and 5 mol% FD fluorene content added, the interface thermal resistance of i‐FD2‐P6M and i‐FD5‐P6M was reduced to 1.0E‐8 and 0.9E‐8 m2 K W−1, respectively. It demonstrated that the introduction of fluorene was effective in reducing the interfacial thermal resistance between the matrix and the filler. When the fluorene content continued to increase to 10 mol%, the thermal conductivity of the nanocomposite started to decrease and the thermal resistance rose to 1.4 E‐8 m2 K W−1. The calculated interface thermal resistance was consistent with the experimental thermal conductivity of the nanocomposites, as shown in Figure 4a,c.4Figurea) Thermal conductivity of i‐FDy‐P6M with different fluorene content; b) thermal conductivity of FDy‐PEEK with different fluorene content; c) interface resistance of i‐FDy‐P6M with different fluorene content; d) tensile strength and elongation at break of i‐FDy‐P6M with different fluorene addition.The mechanical properties of nanocomposites were measured and shown in Figure 4d. The tensile strength of the composites showed a trend of first rising and then falling. When a small amount of fluorene group was introduced, although the crystallinity of the polymer was slightly reduced, the increased bonding force between the resin and the filler would make up for the loss. This is why the tensile strength of i‐FD2‐P6M was enhanced to 106.5 MPa with 2 mol% fluorene groups introduced, which was 10.4% higher than that of i‐P6M. However, continuing to increase the content of fluorene groups would further limit the crystallization ability of the copolymers. The drop in crystallinity led to a significant decrease in the intrinsic strength of the polymer matrix, which cannot even be compensated by the reinforced interfacial bonding. Therefore, the tensile strength of i‐FD5‐P6M started to decrease. The elongation at break of the nanocomposites showed the same trend. With the rising fluorene content, the elongation at break of the nanocomposites increased significantly first and then decreased, reaching the maximum of 8.5% when 2 mol% fluorene was introduced.[34,41–43] A combined consideration of the thermal conductivity and mechanical properties gives that an introduction of 2 mol% fluorene groups is the optimal ratio for our in situ polymerization method to prepare the fluorene‐functionalized PEEK/MWCNT nanocomposites.With the optimal fluorene content determined, the nanocomposites with or without fluorene groups at various MWCNT amounts were selected to be compared. Figure 5a shows the thermal conductivity of i‐PxM without fluorene introduced and i‐FD2‐PxM with 2 mol% fluorene introduced. The thermal conductivity of i‐FD2‐PxM always showed an increment after introducing fluorene and reached 2.41 W m−1 K−1 for i‐FD2‐P10M. Meanwhile, the tensile strength was compared in Figure 5b, showing that the introduction of fluorene allowed the tensile strength of i‐FD2‐P10M to reach 104.3 MPa and the elongation at break increased to 6.4%. In order to elucidate the enhanced the elongation at break of the i‐FD2‐P10M composite, its tensile section was examined using scanning electron microscopy and was presented in Figure 5c,d. As observed from the SEM images, the i‐P10M section exhibited significant agglomeration, leading to stress concentration and a reduction in elongation at break. However, with 2% fluorene group introduced into the polymer chain, the tensile section of the composite became uneven, with the morphology of the resin stretched, and the length of the MWCNT being pulled out of the resin increased significantly. This improvement can be attributed to the significant enhancement in the interfacial bonding between the resin matrix and carbon nanotubes due to the introduction of fluorene groups, leading to better stress transfer to MWCNT during tensile processes and thus improving the strength and elongation at break of the nanocomposites.5Figurea) Thermal conductivity of i‐PxM and i‐FD2‐PxM with different MWCNT additions; b) tensile properties of i‐P10M and i‐FD2‐P10M; c) tensile fracture surface morphology of i‐P10M; and d) i‐FD2‐P10M; e) elevated temperature thermal conductivity of i‐FD2‐P10M; and f) DMA curves of i‐P10 and i‐FD2‐P10M.As high‐temperature resistant materials, the PEEK nanocomposites are usually required to be evaluated at high temperatures. Thermogravimetric analysis was performed on the nanocomposite i‐P10M and i‐FD2‐P10M, and the results are presented in Figure S4 (Supporting Information). It was found that when the fluorene group was incorporated, the composites exhibited thermal stability with a 5% thermal decomposition temperature above 500 °C. This indicates that i‐FD2‐P10M composite can meet the severe processing temperature. The thermal conductivity of i‐FD2‐P10M at elevatored temperatures was measured and shown in Figure 5e. With the increasing temperatures, the thermal conductivity of the nanocomposite gradually rose and reached 2.96 W m−1 K−1 at 150 °C (a temperature close to Tg). Although high temperature would increase the interfacial scattering of phonons, it also improved the intrinsic heat transfer efficiency inside the carbon material.[44] Therefore, the thermal conductivity of the nanocomposite tended to increase with the rising temperature. This showed that the materials still had good performance at high‐temperature conditions. In addition, the DMA measurement was carried out on i‐P10M and i‐FD2‐P10M and their curves were shown in Figure 5f. The prepared nanocomposites maintained high storage modulus (G') before 150 °C. Benefiting from the better interface interaction, G’ of i‐FD2‐P10M was slightly higher than that of i‐P10M. In consistent with DSC test results, i‐FD2‐P10M had a higher glass transition temperature than i‐P10M. To conclude, the introduction of fluorene groups enhanced the interface between the polymer matrix and MWCNT, further limiting the molecular chain movement and improving the heat resistance. After glass transition temperature, G’ of i‐FD2‐P10M was slightly lower than that of i‐P10M, which was mainly affected by the crystallinity of the copolymers.[45]Nanocomposites can also be used as a matrix to prepare multi‐component composites. Following our previous work, a double segregated network composite i‐PEEK/MWCNT@GNS was prepared by replacing the matrix with the fluorene‐functionalized PEEK/MWCNT nanocomposite.[35] As shown in Figure 6, after introducing fluorene groups, the thermal conductivity of i‐FD2‐P10M@20GNS was improved to 3.91 W m−1 K−1, which exhibited a significant increment compared with that of i‐P10M@20GNS. At the same time, the tensile property of the composite was also remarkably improved from 57.7 MPa to 65.8 MPa. The ternary composite's strength enhancement was due to the improved intrinsic strength of the fluorene‐functionalized PEEK/MWCNT nanocomposite matrix. Furthermore, in comparison with other PEEK/fillers thermal conductive composites reported in the literature as shown in Figure 6 and Figure S5 (Supporting Information), the thermal conductivity in our work reached a relatively high level.[13,14,20,46–54] The combination of in situ polymerization and molecular structure design improved the dispersion state of carbon nanotubes, introduced a large amount of MWCNT as a thermally conductive filler, and significantly improved the thermal conductivity interface, resulting in better thermal conductivity of the synthesized nanocomposites. This approach provides a novel view to design multi‐component composites with balanced thermal conductivity and mechanical strength.6Figurea) Comprehensive performance chart of different matrix composite 30 wt% filler; b) comparison of thermal conductivity between the composites in this work and other PEEK/fillers composites reported by literature (this work‐1 represents i‐FD2‐P10M and this work‐2 represents i‐FD2‐P10M@20GNS).To visually demonstrate the heat dissipation capability of the composites, a device was constructed by integrating the composite with an LED light and the heat dissipation effect was tested under the actual working scenario, which was shown in Figure 7a. The real‐time temperature of the composites above the LED was tested by infrared thermography, and the curve was plotted in Figure 7b. The temperature of the material gradually increased with working time, but the surface temperature of i‐FD2‐P10M was consistently lower than the other two references. The final temperature of i‐FD2‐P10M was 17 °C lower than that of pristine PEEK, indicating that i‐FD2‐P10M has superior heat dissipation ability and can significantly reduce the operating temperature of LED devices.7FigureDiagram of a) LED heat sink parts, b) operating temperature curves, and c) infrared thermogram.ConclusionIn the present work, the fluorene‐functionalized PEEK/MWCNT nanocomposites with high thermal conductivity and tensile strength were synthesized through an in situ copolymerization procedure. The introduction of fluorene groups into PEEK can effectively improve the interfacial bonding between the copolymer matrix and MWCNT, thus reducing the interface thermal resistance. The thermal conductivity of the nanocomposite i‐FD2‐P10M with 2 mol% fluorene and 10 wt% MWCNT is increased to 2.41 W m−1 K−1, which is 868% higher than pristine PEEK. Meanwhile, the tensile strength of i‐FD2‐P10M maintain at 104.3 MPa. Based on this nanocomposite matrix, multi‐component composites with segregated network structure are constructed to achieve promising thermal management materials with balanced thermal conductivity and mechanical strength properties. This work provided a generic method to balance the thermal conductivity and mechanical strength of the polymer‐matrix composites and to prepare highly nano‐filled composites with optimized interface adhesion.AcknowledgementsThe authors would like to thank the National and Local Joint Engineering Laboratory for the Synthesis Technology of High‐Performance Polymers for their assistance and use of the facility. This work was financially supported by the Department of Science and Technology of Jilin Province (Grant 20210509034RQ and 20210201117GX) and the Fundamental Research Funds for the Central Universities, the National Natural Science Foundation of China (No. 52003099), the Natural Science Foundation of Jilin Province (YDZJ202101ZYTS002), and the Capital Construction Fund of Jilin Province (2021C039‐1).Conflict of InterestThe authors declare no conflict of interest.Author ContributionsZ.J. and B.J. contributed equally to this work and should be considered co‐first authors. Conceptualization, methodology, investigation, writing original draft. B.J.: Validation, resources, writing original draft. B.Y.: Writing—review and editing, validation. X.L.: Resources, validation. Y.Y.: Investigation, validation. Y.S.: Validation, data curation, writing, reviewing and editing. 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Journal

Advanced Materials Interfaces – Wiley

Published: May 1, 2023

Keywords: carbon nanotubes; interfaces; nanocomposites; thermal conductivity

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