Hydrogen‐Bond‐Mediated Surface Functionalization of Boron Nitride Micro‐Lamellae toward High Thermal Conductive Papers

Wide‐bandgap, layered hexagonal boron nitride (h‐BN) possesses excellent electrical insulation and ultrahigh thermal conductivity simultaneously, offering a perfect candidate for the growing demands of heating dissipations in modern chip industries and power electronics. Hybrids of h‐BN with polymers fulfill the thermal management materials (TMMs) requirement of flexibility, while the composite poses severe challenges in the interfacial bonding and excess thermal resistance. To date, the practical bonding between h‐BN intrinsic surfaces and polymer matrices remains elusive. This work reports on the effective alignment of h‐BN micro‐lamellae by introducing nitrogen‐atoms‐containing polymers as inter‐lamellae bridging mediums. Based on theoretical calculations the hydrogen bonding between polymer chains and the BN surface is revealed by differential charge densities mapping. It is shown experimentally that the neuron‐like polymer bundles strongly bonding surfaces of two neighboring h‐BN platelets as direct, microscopic evidence of the structure models. An extra alignment of h‐BN induced by this strong interfacial interaction leads to a higher degree of h‐BN stacking order, boosting the thermal conduction by eight times. These results reveal one unprecedented method to non‐covalently functionalize the h‐BN surface and expand the TMMs family in the dimension of the filler size, paving the way for exploring the larger‐sized ceramic TMMs.


Introduction
The accumulated heat in the miniature integrated circuits and high-power-density devices aggravates the heat dissipation crisis to the detriment of devices' performance, reliability, and even the safety of human beings when considering applications in, for example, the power battery for electric vehicles. [1,2] The rapid development of contemporary electronics asks for thermal management materials (TMMs) with the ultrahigh thermal conductivity (κ), and most of the time, electrical insulation and mechanical compressibility to better contact chips. [3][4][5][6][7] The composite films with polymers as matrix have been widely investigated and industrially used, with high κ materials acting as fillers, including metal, carbon-based nanomaterials, and insulating ceramic particles. [3][4][5][6][7][8] The light-element hexagonal boron nitride (h-BN), in which boron and nitrogen atoms are arranged sequentially in a honeycomb lattice, [9][10][11] is an excellent candidate of ultrahigh-κ fillers due to its superior electrical insulation (electrical breakdown field: 4-10 MV cm −1 ) and ultrahigh κ (planar: ≈600 W m −1 K −1 ). The most significant challenge for its applications as TMMs is its "chemically inert" surface. [12,13] Unlike graphene flakes which can be easily exfoliated from graphite in solution and simultaneously surface-functionalized by, like the Hummers method, [14][15][16][17] h-BN occupies a higher level of chemical inertness, leading to difficulties in chemical functionalization on the surface, the intercalation of gas molecules or alkali metals into the van der Waals gaps, and sample dispersion in solutions for large-scale exfoliations. [18] To disperse h-BN sheets in solvents for the subsequent assembly with polymer matrices, typically the edge-functionalized strategy was applied. [19][20][21][22][23][24] For example, the organic superacid or ionic liquid has been introduced to physically bond the edge of h-BN via electrostatic interactions, that is, electrondeficient B atoms interacted with cations and anions attached to N atoms by utilizing dangling bonds at edges. [19,20] More readily, by sonicating h-BN materials in pure water, highly negative zeta potentials were revealed for h-BN sheets and attributed to the hydroxyl groups generated at edges (BOH and NOH), benefiting its dispersibility in solutions and the stability of formed dispersion. [21,22] By following this idea, treatments in concentrated alkaline conditions or molten alkalis have been proved to effectively prepare edge-hydroxylated h-BN nanosheets. [23,24] The engineering of edge functionalization desires more edges which can serve as the interacting sites with polymers. Therefore, severe mechanical exfoliations by a long-term probe sonification or intense ball milling were generally applied to reduce the thickness of h-BN flakes (typically, the lateral size as well), producing the h-BN nanosheets with plenty of edges. [25][26][27][28] On a parallel front, much less explored but probably more significant is the non-covalent functionalization of h-BN surface, rather than edges, to provide more bonding sites on the surface for the dispersion and subsequent assembly with polymers. Since the number of atoms at edges is much smaller than that of "center" atoms on surfaces, one efficient surface functionalization would be very helpful and as a novel approach, for h-BN-based TMMs composites. However, up to date, only a few works reported the polymer-BN interaction on surfaces by the π-π stacking-like interaction or Lewis acid-based complex effects but without direct, unambiguous evidence to show whether the bonding sites are actually on h-BN surfaces. [29][30][31][32] The underlying challenge is the missing fundamental understanding of interactions with the chemically inert h-BN surface. [29][30][31][32] In this work, we demonstrate one path to functionalize surfaces of h-BN by using electrostatic interactions. We showcase that the chemically inert but ionically active surface of h-BN can strongly bond with nitrogen-containing polymers, with the supports of theoretical calculations on the single molecule level and unprecedented, distinct experimental images which unambiguously reveal the bonding sites on surfaces. This robust surface linking with polymers leads to enhanced interfacial bridging, [8,33,34] thereby accomplishing the orientational alignment of h-BN micro platelets. We further demonstrate the ultrahigh thermal conduction and mechanical bendability with ultrahigh h-BN contents (up to 89 wt%) as the consequences of the proposed engineering of the surface functionalization.

Surface Functionalization by Electrostatic Interactions
We first theoretically investigated the interaction between an h-BN surface and guanidine with the formula HNC(NH 2 ) 2 , one common nitrogen-containing molecule whose moiety frequently appeared in the main chain of larger organic molecules. Methylguanidine on top of monolayer h-BN was the simplest case to begin with. As shown in Figure 1a, three nitrogen atoms are labeled by N1, N2, and N3, respectively, which are connected to (N1), close to (N2), and far away from (N3) the methyl group (CH 3 ). Note that one hydrogen attached to N2 is protonated, offering one positive charge for the whole molecule, that is, without changing the electrostatic feature of nitrogen cation (N2). We systematically considered three scenarios in which N2, N2 + N3, and N1 + N3 stood on one h-BN surface with three types of upright configurations, named by upright-2, upright-2,3, and upright-1,3, respectively, followed by the structure relaxations (inset of Figure 1a). The corresponding adsorption energies E a are displayed in Figure 1a. We found that in the first configuration the cationic N2 atom exhibited the preferable position on top of the boron atom, due to the electrostatic interaction between N2 and partially, negatively charged boron in h-BN lattices. Moreover, the most energetically favorable configuration (E a = −1.78 eV) was the upright-1,3 one where N1 and N3 atoms were close to N atoms of h-BN lattices. To better understand the molecule-BN interaction, the spatial charge density distribution of the self-standing molecule (D a ), h-BN itself (D b ), and the combination of methylguanidine and h-BN (D c ) were calculated separately. The differential charge density ΔD, that is, D c − (D a + D b ), is shown in Figure 1b. The positive (negative) ΔD indicates the increase (decrease) of the charge density when incorporating the molecule into h-BN. In this way, the newly formed bonding between molecule and h-BN is highlighted by the positive ΔD region, which is further marked by yellow. From Figure 1b one can readily identify the bonding between H atoms (connected to N1 and N3 atoms) in the molecule and N atoms in h-BN. By considering the electronegativity discrepancy between N and H atoms, these two interfacial bonds were recognized as hydrogen bonding, denoted by NH···N. The first N atom was from the molecule, while the latter belonged to h-BN. Furthermore, a more complicated case, bilayer h-BN/dimethylguanidine, gave the same conclusions about molecule/h-BN hydrogen bonds and the energetically preferred upright configuration ( Figure S1, Supporting Information). We also note that the long-chain polymers can provide multiple nitrogen and hydrogen atoms serving as the potential bonding sites with h-BN surfaces, which vastly increase the probability and strength of surface-polymer interactions.
Experimentally, we chose large-dimension h-BN microlamellae, instead of widely-used BN nanosheets, as fillers for two reasons. First, one intuitive idea was that if using largesized h-BN platelets, constructing one heat conduction channel from the source to the sink required fewer building blocks, thereby effectively reducing the number of h-BN/h-BN interfaces and probably interface thermal resistance. And we further substantiated this idea by finite element simulations ( Figure S2, Supporting Information). In the previous studies of thermally www.advmatinterfaces.de Figure 1. Surface functionalization by hydrogen bonding interactions. a) Calculated adsorption energy E a of three configurations for one methylguanidine molecule on a BN monolayer, showing the energetically favored upright-1,3 configuration. b) Differential charge density for the upright-1,3 configuration. The yellow and blue region indicates where the charge density increases and decreases, respectively. One can easily identify the enhanced charge density, or in other words, the bonding, between H atoms connected to atoms N1 and N3 in the molecule and N atoms in h-BN lattices. c) Schematic illustration of one PVA-PHMG-BN film. PHMG simultaneously connects to h-BN flakes and PVA chains by hydrogen bonds, eventually achieving one compact structure of BN-polymer-BN.

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conductive h-BN/polymer composites, the lateral size of h-BN sheets was typically limited to several hundreds of nm, or even smaller, as a result of the compromise with the dispersing challenge in solutions. [19][20][21][22][23][24][25][26][27][28] In other words, few works explored the usage of h-BN micro lamellae. [35] Meanwhile, if the proposed strategy of the surface functionalization worked for largesized flakes, it would also succeed in conventional nanosheets. Second, compared with h-BN nanosheets, micro platelets possessed a much lower ratio of edges and could be easily distinguished from polymer matrices in microscope images. Both are beneficial for achieving and recognizing the basal surface interaction of h-BN with polymers.
As shown in Figure 1c, we thereby introduced polyhexamethylene guanidine (PHMG) to act as a "bridge" between the h-BN and polyvinyl alcohol (PVA) matrix. While PHMG can be anchored on the surface of h-BN through the proposed hydrogen bond interactions, it can simultaneously connect with PVA molecules by hydrogen bonds as well, leading to the strong interactions between h-BN/PVA and then h-BN/ polymer/h-BN in PVA-PHMG-BN films. We further show that, by this inter-lamella bridging, large-sized h-BN platelets can be orientationally aligned and stacked well to form a flexible paper with ultrahigh thermal conductivity.

Bridging Effect of PHMG
The PVA-PHMG-BN films were prepared by the vacuum filtration, while the vacuum filtering itself can roughly align the h-BN lamellae. Precursors with designed loading ratios were applied to the mixture of isopropanol alcohol and deionized water (vol. 1:1), followed by gentle stirring and the mild bath sonication to disperse h-BN lamellae in solvents (see details in Experimental Section). Note that intense exfoliation techniques, such as a ball milling or probe sonication, were not adopted to preserve the size of commercial h-BN powders. As shown in Figure S3, Supporting Information, scanning electron microscope (SEM) images and size statistics of h-BN lamellae exhibit the distribution of 4-14 µm in planar sizes for the as-received samples and that of 2-6 µm after the sonication, guaranteeing in hybrid films the large lateral dimension of h-BN lamellae which was further confirmed by directly observing severalmicron-large flakes in PVA-PHMG-BN films ( Figure S3c, Supporting Information).
The PHMG manifested itself in several aspects. First, compared with the instability of pure h-BN lamellae in solvents with the same mild process (0.37 mg mL −1 , Figure 2a), the introduction of PVA and PHMG essentially improved the solution stability of h-BN lamellae without agglomerations observed after 4 days, even with an ultrahigh loading content of h-BN (up to wt. 80%), preliminarily suggesting the effective decoration of PHMG on surfaces of h-BN platelets. One consequent benefit was that we were able to obtain a series of stable colloidal suspensions with varied loading contents of h-BN (30% to 80%) and the resulting free-standing films by peeling them off from filter membranes (Figure 2b; Figure S4, Supporting Information).
Surprisingly, unexpected flexibilities were revealed for PHMG-mediated composite films. For instance, a film with 40 wt% of h-BN (loading content, notated as PVA-PHMG-BN 40) can be manually folded to be an airplane and restored without apparent breakages (Figure 2b), and when increasing the content of h-BN to an ultrahigh level (80%), the films can still endure bending by being curled into a cylinder (Figure 2b and Figure S4a-d, Supporting Information). However, as control experiments without adding PHMG, a PVA-BN 40 film acted as a mechanically brittle film without any flexibility and was hardly handled by a tweeze which broke the film easily ( Figure 2d). Even worse, the PVA-BN 60 film lost its integrity by cracking into the debris and could not be peeled off as one free-standing film (Figure 2d). Again, the sharp contrasts of the film-forming ability, mechanical strength, and film's bendability were believed as an outcome of the incorporation of PHMG, of which the intercalation between h-BN lamellae increased the h-BN/h-BN bonding and flexibility of films. While bending the nanosheet-based composite films can be realized by the relative slippage and even curling of nanosheets themselves, [26,27] in our experiment the stiffness of 100 nm thick lamellae defies bending by curving each lamella. The variation of inter-lamellae distances upon bending is therefore expected, asking for a strong inter-platelet bonding to restore when retracting the external forces. Previously no works showed the bendability of micro h-BN lamellae-based hybrid films, not to mention with so high h-BN contents (loading ratio: 80 wt%), strongly suggesting the efficient improvement of inter-BN coupling.
Moreover, the microscopic evidence of PHMG bridging effects is even more stunning from cross-sectional SEM images (Figure 2d,e). Without PHMG, the PVA-BN 40 film exhibits one nacre-mimetic structure with h-BN flakes stacking chaotically, displaying a porous structure and the absence of polymers "bonding" on the surfaces of flakes (Figure 2d). In marked contrast, with the incorporation of PHMG the more compact, layer-by-layer morphology is found (Figure 2e). Significantly, the neuron-like filaments, of which two ends are connected to surfaces of two neighboring h-BN lamellae, are routinely observed in every single film we prepared, and we name this bonding "inter-lamellae bridging" (see more images in Figure  S5, Supporting Information), which unambiguously support the microscopic pictures of the PHMG bridging as shown in Figure 1c. We note that to the best of our knowledge, it is the first time to show the microscopic and direct evidence of strong interactions between the h-BN surface (rather than edges) and polymers. [26,31,32] And here, the advantage of using micro h-BN lamellae is emphasized by the capability to clearly distinguish h-BN platelets and polymer bundles in SEM images, while in h-BN nanosheet-based composite films, one hardly tell h-BN itself from images. Considering the same level of vacuum filtration-induced h-BN orientations in two types of films (with and without PHMG), we emphasize the extra alignment of h-BN from PHMG bridging, leading to the compact stacking of lamellae as shown in Figure 2e.
Furthermore, another macroscopic effect of PHMG bridging is the content of polymers after the vacuum filtration. If the polymer matrices did not have strong interactions with fillers, there would be a large discrepancy between the loading and real contents of polymers by pumping the interaction-free polymer out with solvents. As displayed in Figure 2f,g, with the thermogravimetric analysis (TGA) the real mass of h-BN is measured after volatilizing the polymer matrix, including PHMG and www.advmatinterfaces.de PVA, at high temperatures (also see Table S1, Supporting Information). For PVA-PHMG-BN films, a 10-25 wt% increase of h-BN is found when comparing real contents with the loading ones, and inversely proportional to the loading contents of h-BN, for example, 54.9 wt% of actual content for a PVA-PHMG-BN 30 film (an increase of 24.9 wt%) and 89.3 wt% of measured content for a PVA-PHMG-BN 80 film (an increase of 9.3 wt%), suggesting the mass loss of polymer during the process. The presence of polymers in PVA-PHMG-BN films was also confirmed by the energy dispersive X-ray spectroscopy (EDS) and Fourier transform infrared spectroscopy (FTIR) characterizations ( Figure S4e-f, Supporting Information). However, for PVA-BN films without PHMG, both PVA-BN 20 and PVA-BN 40 films exhibit the real contents of h-BN above 90 wt% (more specifically, 94.8 and 95.3 wt%, respectively), demonstrating the poor bonding between PVA and h-BN and the massive polymer loss during the vacuum filtration. Thereby, the as-received PVA-BN hybrid film was almost pure h-BN one, consistent with the mechanical fragileness as shown in Figure 2c.

Enhanced Thermal Conduction by PHMG-Mediated Alignments
To further illustrate the advances of PHMG-induced h-BN alignments, the in-plane and out-of-plane thermal diffusivities (α) of

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PVA-PHMG-BN films were measured by the laser flash method (see the details in Experimental Section), and the thermal conductivity κ of films was calculated by the equation κ = α × ρ × C p , where ρ and C p represent the density and specific heat capacity of the composite films, respectively ( Figure S6, Supporting Information). As shown in Figure 3a, we list κ as a function of real h-BN contents calibrated by TGA measurements. For a PVA-BN film with 95 wt% of the h-BN real content, its in-plane κ is only ≈ 5.8 W m −1 K −1 , which is consistent with its chaotically stacked structure (Figure 2d). With the incorporation of PHMG the κ can reach up to 46.5 W m −1 K −1 , a more than 800% enhancement, again signifying the effective alignment of h-BN fillers by inter-lamellae PHMG bonding. Surprisingly, the in-plane κ of films first increased and then decreased upon h-BN contents, presenting the featured Λ shape with its maximum at an h-BN content of 65.1 wt%. This finding indicates the invalidation of PHMG-induced alignments when having a high h-BN content (>65.1 wt%).
We further performed X-ray diffraction (XRD) characterizations to quantify the degree of stacking order for h-BN micro-lamellae in composite films. As displayed in Figure 3b, the prominent peak is located at 26.8° and assigned to (002) crystallographic plane. And one can find four minor peaks in the range of 40-60°, which can be well indexed as (100), (101), (102), and (004) peaks, respectively. Here, we used the intensity ratio of (100) and (002) peaks (I (100) /I (002) ), that is, the ratio of "vertical" h-BN lamellae (corresponding to (100) peak) and ones along the in-plane direction contributing to (002) peak, as a describer of the degree of the h-BN stacking order. For as-received h-BN powders without any intentional alignment, I (100) /I (002) shows the highest value of 1.71%. Interestingly, this order parameter exhibited the V shape dependence of the h-BN content with a minimum (0.06%) at 65.1 wt%, well correlating with the behavior of the in-plane κ. Moreover, the cross-sectional SEM characterizations (Figure 3c-f) show an apparent alignment variation of h-BN lamellae upon contents. We can easily find that with lower contents h-BN lamellae stack orderly with a well-connected configuration (Figure 3c,d); however, with higher contents (>65.1 wt%), the degree of stacking order decreases with a large number of micro-scale holes (Figure 3e,f), which is also confirmed by densities much lower than the theoretical predictions ( Figure S6b,

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Supporting Information). The PHMG-induced extra alignment of h-BN fundamentally stems from the newly-formed hydrogen bonding and can increase the possibility of one h-BN lamella touching with adjacent another and/or the averaged contact area between h-BN flakes, thereby boosting the thermal conduction.
By combining results of thermal conductivity measurements and stacking order characterizations together, we naturally divided it into two zones as a PHMG-effective zone (Zone I, ≤65.1 wt%) and "not-that-functional" zone (Zone II, >65.1 wt%). In Zone I, apart from the alignment from the vacuum filtration, the PHMG bridging effect worked efficiently, leading to the micro-level, excellent alignment of h-BN platelets, high degree of stacking disorder, and enhanced thermal conductivities by this well-arranged architecture. In Zone II, however, the ratio of h-BN was too high to be aligned by a low-proportion PHMG, displaying the structural voids, chaotic stacking, and depressed thermal conductivities by the discontinuity of transport paths. Note that typically the thermal conduction behavior of one hybrid film largely and positively depends on the content of high-κ fillers. The result shown in Figure 3a seems like counter-intuitive; on the contrary, it reflects another key factor, the degree of h-BN stacking order, which dominates the thermal conduction in some scenarios (here the ultralow proportion of PHMG in Zone II). And this conclusion is further corroborated by the following experiment. However, from the TMMs aspect of view, the high content of fillers (as high as possible) was highly desired, thereby posing one irreconcilable challenge to h-BN orientation managements.

Chemical and Mechanical Dual Alignments of h-BN
Regarding a low level of stacking order in PVA-PHMG-BN films in Zone II, we further introduced the cold pressing as one complementary, post strategy for dually aligning h-BN platelets. As illustrated in Figure 4a, due to the interrupted heat conduction paths by holes and disordered alignments of h-BN, PVA-PHMG-BN films in Zone II experience a poor thermal transport. By applying one uniaxial pressure (along the z direction, ≈20 MPa) for 10-15 mins, benefiting from the micron dimension, h-BN platelets were mechanically forced to be laterally oriented, thereby "squeezing" holes out and increasing the alignments of h-BN for a higher thermal transport efficiency. After pressing, PVA-PHMG-BN films still act as free-standing films with considerable flexibilities (Figure 4b). For example, the compressed PVA-PHMG-BN 80 film (actual h-BN content: 89.3 wt%) can be curled as a cylinder (Figure 4b) and restored completely. To further exam the films' microstructures, we showed cross-sectional SEM images of one PVA-PHMG-BN film before (Figure 4c) and after a cold pressing (Figure 4d). Notably, its thickness was reduced by more than 100% (from ≈45 to ≈20 µm) with much fewer voids. More importantly, the zoomed-in SEM images unambiguously confirmed the pressing-induced alignments of h-BN which supplemented the h-BN orientation control with PHMG bridging. Moreover, I (100) /I (002) was measured by XRD characterizations, exhibiting a significant reduction (from 1.20% to 0.10%, Figure S7, Supporting Information) which was consistent with SEM observations. Parallelly and compactly stacked h-BN lamellae with larger contact areas by our dual alignment process should have an improved heat transport. As exemplified by a PVA-PHMG-BN 80 film which originally had a low thermal conductivity of 10.6 W m −1 K −1 , after the cold pressing, the in-plane κ of this film increased to 66.8 W m −1 K −1 (Figure 4e), a more-than-fivetimes increase when compared with the non-pressed sample and an ≈1150% improvement in comparison of the non-duallyaligned film, substantiating the success of h-BN orientation managements by chemical and mechanical approaches. As shown in Figure 4f, we summarized our results of PVA-PHMG-BN films with those of previously reported h-BN nanosheetsbased composites with polymers, cellulose nanofibrils (CNF), cellulose nanocrystal (CNC), or metals. [23,25,28,[36][37][38][39][40][41][42][43][44][45][46][47] We note that our result exceeds most of h-BN nanosheets-based hybrid films significantly and shows flexibility even at an ultrahigh h-BN content, emphasizing the strong fillers/polymer and consequent fillers/fillers bonding stemming from PHMG and dual orientational alignments.
For direct examinations of the thermal transport behavior in real situations, the pure PVA film, PVA-PHMG-BN 60 films without and with cold pressing were selected as heat spreaders with a thermostatic hot-plane heat source. Practical surface temperatures of films were monitored simultaneously by a thermal infrared imager ( Figure S8, Supporting Information). As shown in Figure 4g, after 90 s, surface saturation temperatures of three films are 63, 57.2, and 49.9 °C, respectively, showing superior heat dissipation of the dual-aligned PVA-PHMG-BN film. As a result, the flexibility and high efficiency of our electrically insulating heat spreaders could provide broad prospects for thermal management applications in electronic packaging and printed flexible circuit boards.

Conclusion
Pursuing TMMs with the ultrahigh κ is one of the ultimate answers to the heat dissipation crisis. Apart from the exploration of novel TMMs, the managements of fillers' architectures, such as a 3D continuous skeleton and/or highly oriented stacking, are indispensable alternatives. Engineering the assembly level at the microscopic scale is actually on the basis of a fundamental understanding of interactions between fillers and polymer matrices. Taking h-BN as one example, previously the missing of effective surface functionalization of h-BN denies the usage of h-BN micro lamellae, and the community has to shrink down the filler's sizes to obtain sufficient dispersibility and bonding with polymer matrices (typically at edges). [17,[19][20][21][22][23][24][25][26][27] The strategy of the interfacial hydrogen bonding with nitrogen-containing polymers expands the library of high-κ fillers in one new dimension, that is, the size dimension, and would probably work for other TMMs systems, further paving the way to exploring the larger-sized ceramics as fillers. Moreover, we expect an advanced performance if using h-BN micro-lamellae and nanosheets together. Since, in principle, the PHMG bridging works for h-BN nanosheets as well, sub-micron voids in current PVA-PHMG-BN films can be compensated with nanosheets, giving rise to improved thermal transport efficiency.

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In summary, we have demonstrated one unprecedented surface functionalization of h-BN, in which the hydrogen bonding between nitrogen-containing polymer and h-BN basal surface was utilized. We first corroborated this type of non-covalent interactions theoretically and further verified the PHMG bridging effect in micro-lamellae-based composite films from multiple aspects, including the stability of the dispersion, the unexpected bendability even at the ultrahigh content of h-BN (89.3 wt%), microscopic polymer bundles bonded to h-BN surfaces, and the considerable polymer contents after vacuum filtering. This strong polymer-filler interaction led to a high degree of an h-BN stacking order and high thermal conduction. Combined with a post mechanical pressing, dual alignments of h-BN gave rise to highly-ordered h-BN stacking and consequent ultrahigh thermal conductivity, an ≈1150% increase when compared with untreated samples. Our findings promote a fundamental understanding of non-covalent h-BN/polymer interactions and practical exploration of ultrahigh thermal composite films with large-sized fillers, opening a door for new types of high-performance TMMs. [12,13]
Preparation of h-BN-Based Composite Films: Isopropyl alcohol and deionized water were mixed with 1:1 weight ratio, and typically 0.6 g of commercial BN micropowder (ENO MATERIAL, 5-10 µm) was added to the solution, followed by stirring for 10-20 min (600 rpm) to obtain the pre-mixtures. The bath ultrasonication was applied for 12-15 h. After 36 h, the supernatant of h-BN dispersion was collected. For PVA-BN film, PVA was first added into deionized water and stirred at 95 °C until the solution was clear. The PVA solution was then mixed with the h-BN dispersion with the designed ratio, and stirred for 8 h. For PVA-PHMG-BN films, PHMG was first dissolved in deionized water (10 mg mL −1 ) and added into the h-BN dispersion with the designed ratio, followed by the bath ultrasonic treatment for 25 min and stirring for 90 min (700 rpm). The PVA was added into deionized water (30 mg mL −1 ) and stirred at 95 °C (600 rpm) until the solution is clear. Then PVA solution was mixed with PHMG-BN mixtures with further stirring at 40 °C for 10 h. For all of composite films, the total mass of h-BN and polymers was kept at 80 mg. The composite films were produced by standard vacuum-assisted filtration with PES membranes (pore size: 0.22 µm). After the filtration, the products were naturally dried in air, and then the composite films were mechanically peeled off from PES membranes.
Characterizations: The microscopic morphologies of commercial h-BN source, PVA-BN, and PVA-PHMG-BN films were observed by a field emission scanning electron microscope (FE-SEM, Hitachi, S4800). The crystal qualities and lattice vibrations were determined by the SmartRaman confocal-micro-Raman module (100×, NA = 0.9 of objective lens) coupled with a Horiba iHR320 spectrometer and a charge-coupled device (CCD) detector. FTIR spectra were obtained using BRUKER, VKRTEX-70. X-ray diffractions characterizations were done by Rigaku, SmartLab 9KW. TGA was obtained by SDT Q600 (TA). The thermal diffusivities of films were measured by the laser flash analyzer (NETZSCH). The specific heat capacities of the films were measured by the differential scanning calorimetry (Q2000, TA).

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.