# Fine-tuning of core-shell 1D nanoparticles for thermally conductive, yet electrically insulating, 3D-printable polymer nanocomposites

Antoine Bodin, Thomas Pietri, Caroline Celle, Jean-Pierre Simonato

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Antoine Bodin, Thomas Pietri, Caroline Celle, Jean-Pierre Simonato. Fine-tuning of core-shell 1D nanoparticles for thermally conductive, yet electrically insulating, 3D-printable polymer nanocomposites. Materials Chemistry Frontiers, 2024, 8, pp.3949-3961. 10.1039/D4QM00520A. cea-04850128

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This article can be cited before page numbers have been issued, to do this please use: A. Bodin, T. Pietri, C. Celle and J. Simonato, *Mater. Chem. Front.*, 2024, DOI: 10.1039/D4QM00520A.

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# Fine-tuning of core-shell 1D nanoparticles for thermally conductive, yet electrically insulating, 3D-printable polymer nanocomposites

Antoine Bodin, Thomas Pietri, Caroline Celle, Jean-Pierre Simonato*

Univ. Grenoble Alpes, CEA, LITEN, F-38000 Grenoble, France

E-mail: jean-pierre.simonato@cea.fr

## ABSTRACT

Core-shell nanostructures are particularly interesting for the development of dual-property nanofillers for nanocomposites. In this study, advanced materials compatible with the commonly used Fused Deposition Modeling (FDM) 3D printing technique are reported for heat dissipation applications. Core-shell nanowires based on a highly thermally conductive silver core coated with an electrically insulating silica shell are investigated. The heat dissipation performance of polycarbonate nanocomposites is analyzed using a comprehensive set of thermal, electrical, mechanical, and rheological characterizations to determine the optimal silica nanolayer thickness. We demonstrate that these core-shell nanofillers give access to both high thermal conductivity up to $2.08 \pm 0.05 \text{ W m}^{-1} \text{ K}^{-1}$, and electrically insulating behavior (electrical resistivity $> 10^{12} \Omega \text{ cm}$) at only 3 vol% loading, while retaining very good mechanical strength. The high dispersion and interfacial cohesion of the nanomaterials with the matrix play a key role in achieving these performances. Moreover, thanks to the alignment of the 1D nanofillers during the FDM printing process, the thermal conductivity of the PC nanocomposite reaches an unprecedented value of $3.48 \pm 0.06 \text{ W m}^{-1} \text{ K}^{-1}$ in the printing direction, i.e. a fifteen-fold increase over the thermal conductivity of neat PC.

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Introduction

The rapid growth of microelectronics and energy storage technologies is giving rise to increasingly powerful systems that simultaneously require greater miniaturization and further integration of components, thereby generating undesirable and potentially damaging heat for systems in operation.¹,² Indeed, novel fast-charging batteries with high energy densities triggers high heat generation rates during charging and discharging cycles. This heat accumulation can lead to thermal runaways threatening the integrity of batteries, raising safety issues.³⁻⁵ Similarly, the development of increasingly powerful microprocessors in ever-shrinking dedicated volumes leads to the build-up of hotspots since high integration of transistors has resulted in the escalation of power dissipation and heat generation in microelectronic devices. If heat is not properly dissipated, overheating of the electronic components will inevitably have a detrimental influence on their lifespan and reliability.⁶,⁷ The lack of space for convection cooling systems in portable devices entails the need for developing a new class of heat dissipative structural materials. Henceforth, improving the performance of batteries and electronic systems now heavily depends on the efficiency of their thermal management. In addition to the good thermal conductivity required to facilitate heat dissipation, these new materials also need to provide good electrical insulation to avoid electrical issues such as short-circuits, power overconsumption or signal propagation delay.⁷⁻⁹

Thermoplastic polymer materials offer great potential to meet these needs, thanks to their intrinsic electrical insulation, low cost and ease of processing. However, pure polymeric materials generally exhibit very limited thermal conductivity (< 0.3 W m⁻¹ K⁻¹).¹⁰ A promising strategy for improving thermal conductivity while preserving electrical insulation of polymer-based materials is to incorporate thermally conductive yet electrically insulating fillers to the polymer matrix. Several types of thermally conductive fillers have been studied over the past few years including particles made of carbon,¹¹⁻¹⁴ metals,¹⁵,¹⁶ and ceramics.¹⁷⁻¹⁹ Carbon-based and metallic fillers generally show very high intrinsic thermal conductivity. However, they also exhibit high electrical conductivity.¹⁴,²⁰⁻²² Ceramics lack free electrons, they are therefore inherently electrically insulating, but corollary less thermally conductive.¹⁰,²³ In order to take advantage of

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DOI: 10.1039/D4QM00520A

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the high thermal conductivity of electrically conductive fillers, core@shell nanostructures combining a highly thermally conductive core and an electrically insulating shell have been investigated.²⁴⁻²⁸ In order to prevent electrical conduction throughout the composite, the shell must efficiently hinder the transport of electrical charges between nanofillers. Shells made of ceramic material such as silica (SiO₂) appear to be ideal candidates for meeting this requirement.²⁹⁻³³ Compared with larger particles, nanofillers are able to build many more thermally conductive paths through the composite, which is beneficial for thermal dissipation purposes.³⁴ Although much less studied than 0D nanoparticles, 1D nanofillers with high aspect ratios are preferable as they ensure improved thermal conductivity through the formation of a 3D percolating network at lower loading ratios, thereby avoiding high viscosities and processability issues ³⁵⁻³⁷. This is particularly important in our study since we aim at developing nanocomposites printable by the FDM 3D printing technology, which adds an extra level of difficulty regarding the necessary achievement of complementary properties to reach good printability (e.g. adequate rheology, thermal stability, etc.).³⁸ The possibility of using 3D printing is a major expected outcome of this study, the aim being to produce complex shapes designed on demand to optimize the thermal dissipation of various devices.³⁹

In the present study, we report on the use of silver nanowires (AgNWs) as the core material for core@shell nanostructures, as silver is the metal displaying the highest thermal conductivity, and because some silver-based core-shell derivatives have already displayed interesting performances for non-printable nanocomposites.³²,⁴⁰,⁴¹ On another note, AgNW@SiO₂ core-shell nanostructures have also been used recently in other fields of research, showing that they might be of interest in a wide range of applications.⁴²,⁴³ We demonstrate how a fine tuning of conformal silica nanolayers on the surface of AgNWs induces major changes in the performances of polycarbonate (PC) nanocomposites. The effect of the silica interface with the PC matrix and the influence of its thickness on thermal, electrical, rheological and mechanical properties of the nanocomposites are presented. In addition, the influence of the 3D printing process on the orientation and alignment of the 1D nanostructures was investigated and the resulting variations in thermal conductivity are discussed.

This article is licensed under a Creative Commons Attribution 3.0 Unported Licence.

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Results and Discussion

AgNW@SiO₂ core-shell nanofillers

The thickness of the SiO₂ layer on AgNWs was controlled through a fine adjustment of the TEOS content during the synthesis (Fig. 1a). TEOS amounts of 0.325 %, 0.65 %, 1.30 %, 2.00 % and 4.00 % respectively yielded SiO₂ thicknesses of 5, 11, 20, 27 and 43 nm on the surface of AgNWs, leading to AgNW@SiO₂ core-shell samples referred hereinafter as AgNW@SiO₂-5, AgNW@SiO₂-11, AgNW@SiO₂-20, AgNW@SiO₂-27 and AgNW@SiO₂-43 respectively. FE-SEM images of pure and silica coated AgNWs are displayed in Fig. 1b-h.

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<figure><img src="image_7.png" /><figcaption>**Fig. 1.** (a) Synthesis process of AgNW@SiO₂ nanostructures; FE-SEM images of (b) pure AgNWs and (c-g) AgNW@SiO₂ core-shell nanowires with thicknesses varying from 5 to 43 nm (5 kV accelerating voltage). Insets are images taken at 10 kV accelerating voltage; (h)</figcaption></figure>

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photomontage of an AgNW core coated with a SiO₂ shell of increasing thicknesses (10 kV Article Online DOI: 10.1039/D4QM00520A accelerating voltage); (i) EDX mapping of a single AgNW@SiO₂-27 nanowire.

Applying an accelerating voltage of 10 kV revealed the SiO₂ shell through contrast improvement between the heavy Ag core and the lighter SiO₂ shell. This is highlighted in **Fig. 1b-g** insets and in the high-resolution images of lateral views of each synthesized AgNW@SiO₂ core-shell nanostructure presented in **Fig. 1h**. All samples were coated by a smooth and conformal SiO₂ layer, except for AgNW@SiO₂-5 (**Fig. 1c**) which displayed a layer mainly made of agglomerated and unmerged SiO₂ nanoparticles, and AgNW@SiO₂-11 (**Fig. 1d**) which exhibited a slightly rougher SiO₂ layer compared to other samples. The EDX analysis performed on AgNWs treated with TEOS (**Fig. 1i**) revealed the Ag core as well as both Si and O emission lines (**Fig. S2**), confirming the existence of a continuous SiO₂ shell at the surface of AgNW@SiO₂-20, AgNW@SiO₂-27 and AgNW@SiO₂-43. A set of complementary analyses using FTIR spectroscopy and XPS was carried out to confirm unambiguously the chemical composition and the core-shell structure of the nanomaterials (**Fig. S3** and **Fig. S4**). Contrary to what had been observed with pure AgNWs that irreversibly agglomerate once filtrated, all AgNW@SiO₂ samples were easily redispersed in ethanol or chloroform after vacuum filtration (**Fig. S5**), proving that the SiO₂ shell improves dramatically the dispersion of the AgNWs.

The condensation of TEOS on the AgNWs was made possible thanks to a thin residual layer of polyvinylpyrrolidone (PVP) adsorbed on the metallic surface after the synthesis of AgNWs (highlighted by the carbon emission line spotted in EDX mappings presented in **Fig. S2**). Indeed, the polyol synthesis of AgNWs requires the use of PVP which acts notably as a very thin encapsulating layer.44,45 During the shell fabrication, PVP functional groups assist the SiO₂ layer synthesis, acting as preferential condensation sites for grafting hydrolyzed TEOS molecules onto the PVP enolized form.46 The SiO₂ nanolayer formation can be described by a nucleation/growth mechanism as depicted in **Fig. S6**. At the very beginning of the SiO₂ nanolayer synthesis, hydrolyzed TEOS molecules (Si(OEt)₄·(OH)ₙ with 1 ≤ n ≤ 4) start to condense on the AgNW surface through a heterogeneous nucleation mechanism. Following this preliminary nucleation

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step, the more hydrolyzed molecules reach the nanowire surface, the more the reaction shifts towards a condensation and coalescence mechanism until a continuous SiO₂ nanolayer is formed.

The initial concentration of TEOS for the formation of a conformal and continuous SiO₂ nanolayer lies between 0.65 and 1.30 % v/v to ethanol (corresponding to AgNW@SiO₂-11 and AgNW@SiO₂-20 respectively). During the formation of the silica layer, pure silica particles can form. In order to purify the core-shell nanowires, the solution is vacuum filtered several times over a 1.2 µm PTFE membrane to get rid of free silica nanoparticles. Once purified, the AgNW@SiO₂ core-shell nanostructures with various SiO₂ shell thicknesses were dispersed in the PC matrix to produce PC-AgNW@SiO₂ nanocomposites.

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## Morphology of PC-AgNWs and PC-AgNW@SiO₂ nanocomposites

Pure AgNWs and the five AgNW@SiO₂ samples were infused within the PC matrix to investigate the influence of the SiO₂ shell thickness on the physical properties of the nanocomposites. PC was chosen because of its very interesting mechanical properties as well as its superior dimensional and thermal stabilities. Moreover, PC is a high-performance, heat-resistant thermoplastic polymer that can be easily processed by FDM 3D printing ⁴⁷. Before studying the PC-AgNW@SiO₂ nanocomposite, loadings of 0.5, 1, 3, 5 and 7 vol% were studied for the PC-AgNW reference composite, which permitted to choose the most suitable loading ratio (3 vol%) for the following PC-AgNW@SiO₂ study (see supplementary information for details).

Fractured surfaces of PC-AgNWs and PC-AgNW@SiO₂-5 to PC-AgNW@SiO₂-43 samples loaded with 3 vol% of nanofillers were observed by FE-SEM to appraise the dispersion of the nanofillers and their interfacial compatibility with the PC matrix (Fig. 2b-d). The micrograph of the PC-AgNWs sample (Fig. 2b) clearly shows a separation between the PC matrix and the AgNWs which are mainly agglomerated. Moreover, it was observed that pure AgNWs do not exhibit high compatibility with the PC matrix (Fig. 2c). The repulsive effect encountered between the nanofillers and the polymer matrix leads to low adhesion and creates voids between AgNWs and PC, which results in an increase of the interfacial thermal contact resistance. On the contrary, the samples with AgNW@SiO₂ nanofillers display a much better filler-matrix

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interaction as individual core-shell nanowires are fully embedded and well distributed within the nanocomposites (Fig. 2d-e). This behavior has already been reported for epoxy nanocomposites, where the SiO₂ nanolayer coated on the AgNWs were shown to enhance the wetting of epoxy on the AgNW@SiO₂ surfaces, due to the presence of hydroxyl groups on the silica surface.³² The AgNW@SiO₂ nanofillers were homogeneously and highly dispersed in the PC matrix as a result of a better interfacial cohesion, directly through the solution mixing process. No significant difference was observed on fractured surfaces between the various PC-AgNW@SiO₂ nanocomposites (Fig. S8). As expected, these images of nanocomposite fractured surfaces show that nanofillers are oriented in a specific plane, in a direction perpendicular to the hot-pressing axis. Additionally, adding AgNWs or AgNW@SiO₂ nanofillers to the PC matrix did not affect its thermal stability, as observed from TGA curves displayed in Fig. S9. The onset of thermal degradation was around 400 °C for all samples and the residual mass increased slightly as the SiO₂ shell thickness increased.

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<figure><img src="image_12.png" /><figcaption>Fig. 2. (a) Synthesis process of PC-AgNW@SiO₂ nanocomposites. FE-SEM images of fractured</figcaption></figure>

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surfaces of (b,c) PC-AgNW and (d,e) PC-AgNW@SiO₂ bulk samples. (f) Illustration showing the influence of the SiO₂ coating quality and thickness on electron and phonon flows at the interface of nanofillers embedded in the PC matrix. (g) Through-plane and in-plane thermal conductivity of bulk, hot-pressed samples at 25 °C and (h) thermal conductivity anisotropy ratio with an inset showing the orientation of the nanofillers in the plane perpendicular to the hot-pressing axis. (i) Electrical resistivity of PC nanocomposites filled with pure AgNW and AgNW@SiO₂ nanowires.

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## Thermal conductivity of bulk, hot-pressed nanocomposites

Through-plane and in-plane thermal conductivity values of hot-pressed samples made of neat PC and nanocomposites are displayed in **Fig. 2g**. Adding thermally conductive AgNWs to the PC matrix greatly improved the thermal conductivity of the resulting nanocomposite. Raw AgNWs increased the thermal conductivity in both through-plane and in-plane directions by 82 % (up to 0.42 ± 0.01 W m⁻¹ K⁻¹) and 213 % (up to 0.73 ± 0.02 W m⁻¹ K⁻¹) respectively. Silica-coated AgNWs further increased the thermal conductivity up to 0.66 ± 0.03 W m⁻¹ K⁻¹ and 2.08 ± 0.05 W m⁻¹ K⁻¹ for PC-AgNW@SiO₂-20, corresponding to thermal conductivity enhancements of 187 % and 800 % over neat PC and 60 % and 190 % over PC-AgNW, for through-plane and in-plane measurements respectively (see also **Table S1**).

Such differences in thermal conductivity performances between through-plane and in-plane measurements are explained by one of the special feature of the hot-press shaping process, which tends to orient 1D nanofillers in the perpendicular plane to the pressing direction. This induces higher in-plane thermal conductivity values and high thermal conductivity anisotropy ratios (**Fig. 2h**). The thermal conductivity results obtained in this work stand out from those found in the literature concerning thermally conductive and electrically insulating core-shell and ceramic fillers, as illustrated by the graph exhibited in **Fig. S10**.

The large increase in thermal conductivity for PC-AgNW@SiO₂ samples can be ascribed to the better dispersion of AgNW@SiO₂ core-shell nanowires and the enhanced interfacial
compatibility with the PC matrix (Fig. 2d-e). The combined effects of these two phenomena create a high-performance thermally conductive network deeply anchored in the PC matrix. It is also supposed that the SiO₂ shell acts as an intermediate layer between the hard silver core and the soft PC matrix that alleviates the modulus mismatch at the filler-matrix interface. A high modulus mismatch is indeed responsible for high boundary thermal resistance and phonon scattering caused by vibration frequency incompatibilities existing between two atomically different phases (in this case the metallic AgNWs and the organic PC matrix). Since the thin SiO₂ ceramic shell has a modulus value falling between those of AgNWs and the polymer matrix,³² it creates an easy-to-cross bridge for phonons at the interface, thus reducing their scattering²⁹,³²,⁴⁸ (Fig. 2f). An optimal SiO₂ thickness corresponding to the highest measured thermal conductivity has been determined around 20 nm (PC-AgNW@SiO₂-20). For thinner thicknesses, the lower thermal conductivity is certainly due to an inferior coating quality as observed in FE-SEM images (Fig. 1c, d), detrimental to an efficient phonon transfer at the filler-matrix interface (Fig. S8b', c'). It seems reasonable to assume that the decrease in thermal conductivity values measured on samples with a shell thickness above 20 nm can be explained by the combination of two factors. The first possible reason is associated with the low intrinsic thermal conductivity of the amorphous SiO₂ layer (≈ 1.5 W m⁻¹ K⁻¹),⁴⁹ gradually reducing the group velocity of phonons throughout the interface (Fig. 2f). This decrease and reversing trend in thermal conductivity enhancement has already been observed in other different nanocomposite systems using thick SiO₂ shells in core-shell nanostructured nanofillers.³³,⁴⁶,⁵⁰,⁵¹ The second possible explanation for this drop in thermal conductivity at higher SiO₂ shell thicknesses is associated with a further increase of modulus mismatch. Indeed, previous studies have shown that the elastic modulus of thin film materials depends on the film thickness.⁵²⁻⁵⁴ At low thicknesses (a few nanometers), the elastic modulus of the material decreases with thickness, which alleviates the modulus mismatch. On the opposite, thicker layers gradually get closer to the modulus of the bulk material, as thickness increases. This rationale has already been experimentally verified by two research groups studying the effect of a TiO₂ coating on AgNW⁵² and CuNW²⁸. Jiang *et al.*⁵² found that a 30 nm thick TiO₂ coating on AgNWs had a detrimental effect on thermal conductivity of an

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Materials Chemistry Frontiers Accepted Manuscript

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epoxy nanocomposite, resulting in lower performance than that obtained with pure AgNWs. By contrast, in the study of Ahn et al.,²⁸ the TiO₂ coating on CuNWs was ten times thinner. Under these circumstances, the modulus mismatch between the polymer matrix and CuNWs was lower and led to an increase of the thermal conductivity. The results obtained in the present study suggest that the low intrinsic thermal conductivity of SiO₂ and the increasing modulus mismatch associated with an increasing layer thickness would outweigh the good dispersion and the enhanced filler-matrix cohesion, reversing the trend in the thermal conductivity improvement. The temperature dependence of through-plane and in-plane thermal conductivities of the PC nanocomposites are shown in Fig. S11. As expected, thermal conductivity values rise slightly in a linear fashion from 25 °C to 125 °C, up to 2.29 ± 0.04 W m⁻¹ K⁻¹ for PC-AgNW@SiO₂-20 at 125 °C, which is attributed to the gradual increase in the specific heat capacity of the nanocomposite with temperature.

## Electrical resistivity of PC nanocomposites

In addition to good thermal conductivity, the other essential property sought for PC nanocomposites is efficient electrical insulation. Any material is considered electrically insulating if its electrical resistivity surpasses 10⁹ Ω cm.⁷ Electrical resistivity values are displayed in Fig. 2i. Not surprisingly, neat PC exhibited the highest electrical resistivity (1.02 ± 0.41 x 10¹⁵ Ω cm). However, adding only 3 vol% of pure AgNWs to the PC matrix boosted its electrical conductivity by almost 10 decades (the resistivity plummeted to 4.74 ± 3.40 x 10⁵ Ω cm). On the other hand, the more the SiO₂ coating grows on the AgNWs surface, the more the PC-AgNW@SiO₂ nanocomposite becomes resistive. More specifically, for PC-AgNW@SiO₂-5 and PC-AgNW@SiO₂-11 where the coatings did not cover the entire surface of AgNWs, leaving some gaps as illustrated in Fig. 2f, it was found that electrical resistivity was not satisfactory. It was only at thicknesses of 20 nm and above that the SiO₂ coating was fully effective (PC-AgNW@SiO₂-20 displayed an electrical resistivity of 1.13 ± 0.68 x 10¹² Ω cm), providing complete and satisfying electrical insulation (Fig. 2f, i). Complementary electrical characterizations including the measurements of the dielectric constant, dielectric loss and

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electrical conductivity as a function of frequency were performed (Fig. S12a, b). Breakdown voltage characterization was also investigated (Fig. S12c, d) and confirmed the beneficial effect of the SiO₂ shell on electrical performances of PC-AgNW@SiO₂ nanocomposites, which is in line with expectations for the nanostructures used in this work.

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## Thermal imaging

To assess the heat dissipation performance of PC nanocomposites, all samples, including neat PC, were subjected to a heating and a cooling cycle while their surface temperature was recorded live by an infrared thermal imaging camera. Thermal images for heating and cooling processes are displayed in Fig. 3a and Fig. 3b respectively. Corresponding curves of surface temperature variation vs. time are presented in Fig. 3c and Fig. 3d. PC nanocomposites exhibited significantly faster heating and cooling kinetics than neat PC, as highlighted by the calculated slope coefficients on each heating and cooling curves (see insets in Fig. 3c, d). This observation was expected given the thermal conductivity results obtained previously. Interestingly, the calculated slope coefficient values on the heating and cooling curves followed the same trend as the thermal conductivity results, highlighting an optimum value for PC-AgNW@SiO₂-20. Among the different nanocomposites, those filled with AgNW@SiO₂ core-shell nanostructures were the fastest to heat up and to cool down, demonstrating once again their ability to dissipate heat more efficiently. Precisely, PC-AgNW@SiO₂-20 was the sample exhibiting the best heating and cooling performances. After heating for 30 s, it already reached 76 °C while PC-AgNW was only at 68 °C and neat PC at 55 °C (Fig. 3c). The temperature of the same sample dropped down to 29 °C after only 20 s of cooling from an initial surface temperature of 80 °C, whereas PC-AgNW was at 35 °C and neat PC at 43 °C (Fig. 3d). Thermal cycling performed on the edge of the disk-shaped samples revealed the existence of a pronounced thermal gradient in the in-plane direction for neat PC (Fig. 3c). The temperature distribution for PC-AgNW@SiO₂ nanocomposites was much better compared to other samples. The model depicted in Fig. 3f illustrates this observation through the formation of a high-performance thermally conductive network achieved by well-dispersed AgNW@SiO₂ nanofillers deeply anchored in the PC matrix. Therefore, PC

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nanocomposites with AgNW@SiO₂ nanowires, and especially the one using AgNW@SiO₂. This Article Online DOI: 10.1039/D4QM00520A exhibit better thermal exchange kinetics, resulting in superior cooling abilities.

<figure><img src="image_17.png" /><figcaption>**Fig. 3.** Infrared thermal images of neat PC and PC nanocomposites during (a) heating and (b) cooling cycles; temperature variation over time during (c) heating and (d) cooling cycles. The insets in (c) and (d) correspond to the calculated slope coefficients for heating and cooling cycles respectively; (e) infrared thermal images of neat PC and PC nanocomposites during heating on the edge at t + 90 s and (f) illustration showing the construction of a high-performance thermally conductive network throughout the PC matrix using AgNW@SiO₂ nanofillers.</figcaption></figure>

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PC matrix, tan δ values dropped close to zero as G' overpassed G" over the entire frequency range, indicating a predominant elastic behavior for PC nanocomposites. More importantly, the rheological properties of PC-AgNW@SiO₂-20 and in particular its viscosity (blue star data points in Fig. 4c) were compatible with FDM 3D printer requirements.

## Tensile properties

We chose to assess the mechanical properties of the nanocomposites using tensile tests on neat PC, PC-AgNW and PC-AgNW@SiO₂-20 loaded at 3 vol%, since PC-AgNW@SiO₂-20 achieved the best thermal conductivity results. Tensile measurement results are displayed as averaged stress-strain curves in Fig. 4d. PC nanocomposites exhibited higher Young's modulus but lower tensile strengths and elongations at break than neat PC, as reported in Table S2. The addition of AgNWs to the PC matrix increases its stiffness but also weakens the overall structure. However, AgNW@SiO₂-20 core-shell nanowires lead to better tensile strength and elongation at break than pure AgNWs, and even better Young's modulus. Consequently, PC-AgNW@SiO₂-20 nanocomposites are stiffer than PC-AgNW, but also stronger and tougher. This is again a direct consequence of the enhanced filler-matrix interfacial adhesion while using silica-coated AgNW nanofillers. Mechanical specifications are obviously application dependent, satisfyingly the obtained values reported in Table S2 meet the expected ones for battery casing which was the first targeted application.

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<figure><img src="image_20.png" /><figcaption>**Fig. 4.** (a-c) Rheological and (d) tensile properties of neat PC, PC-AgNW and PC-AgNW@SiO₂ nanocomposites: variations of (a) storage modulus, (b) loss modulus and (c) complex viscosity as a function of frequency at 250 °C; (d) averaged stress-strain curves of neat PC, PC-AgNW and PC-AgNW@SiO₂-20 nanocomposites.</figcaption></figure>

### 3D printing of AgNW@SiO₂-20 nanocomposites

1D nanofillers are known to give anisotropic properties due to their high aspect ratio. To make the most of this anisotropy, the nanofillers can be oriented into a certain direction to maximize the thermal conductivity. The FDM process has become one of the most popular 3D printing technique owing to its easy set-up and low cost. One of the distinctive features of 3D printing of polymer nanocomposites using FDM is that nanofillers tend to align in the printing direction ⁵⁹,⁶⁰. 3D printed test specimens are presented in **Fig. S14**. Neat PC and PC-AgNW@SiO₂-20 specimens were both printed in 0°/0° unidirectional and 0°/90° criss-crossed

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successive layers (Fig. 5a) through a 0.4 mm heated nozzle moving in the (x,y) plane (Fig. 5b) [1] (new Article Online DOI: 10.1039/D4QM00520A). Several cuts were performed on each disk-shaped punched sample to analyze the process-related microstructural arrangement of the nanocomposite, in both unidirectional and criss-crossed printing directions (Fig. 5a, c). The different cutting styles are schematically presented in Fig. 5c.

For the 0°/0° unidirectional print, one cut parallel (1) and another perpendicular (2) to the printing direction were performed (Fig. 5c₁-c₂). For the 0°/90° criss-crossed print, a normal cut (3) and a 45° cut (4) were done (Fig. 5c₃-c₄). Associated FE-SEM images are presented in Fig. 5d₁-d₄.

Insets show the stacking orientation of the successive 3D printed layers for each cut through lower magnification images. The high-magnification FE-SEM images clearly show the orientation and alignment of all AgNW@SiO₂-20 nanofillers in the same direction which corresponds to the printing direction (represented as yellow arrows or arrowheads in Fig. 5d). The thermal conductivity properties associated with this orientation were then measured through-plane and in-plane for each sample. The results are presented in Fig. 5e and compared with bulk properties.

The thermal conductivity of neat PC was not affected by the 3D printing process, neither through-plane nor in-plane directions. Interestingly, the through-plane thermal conductivities (z direction) of PC-AgNW@SiO₂-20 decreased in both unidirectional and criss-crossed printing directions compared to the hot-pressed bulk sample. This can be explained by the fact that almost all nanofillers were oriented in the (x,y) plane, which is perpendicular to the through-plane thermal diffusivity testing direction (z), either in the unidirectional or in the criss-crossed sample. This was not the case for bulk samples, where some nanofillers remained oriented in the testing direction after hot-pressing (random orientation).

Materials Chemistry Frontiers Accepted Manuscript

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<figure><img src="image_24.png" /><figcaption>**Fig. 5.** (a) Printing patterns for test specimen filled with a rectilinear infill; (b) picture of the print head and the printing nozzle used in this work; (c) cross-sectional cutting styles: (c₁) 0/0 parallel cut, (c₂) 0/0 perpendicular cut, (c₃) 0/90 normal cut and (c₄) 0/90 45° cut; (d) corresponding FE-SEM images of (d₁) 0/0 parallel cut, (d₂) 0/0 perpendicular cut, (d₃) 0/90 normal cut and (d₄) 0/90 45° cut. The yellow arrows and arrowheads indicate the printing direction. (e) Through-plane and in-plane thermal conductivity along with corresponding anisotropy ratios of bulk and 3D printed neat PC and PC-AgNW@SiO₂-20 nanocomposite as a function of the printing raster angle and</figcaption></figure>
in-plane cutting style; (f) 3D printed demo parts of neat PC and PC-AgNW@SiO₂-20 with corresponding thermal images during the heating cycle (150 s in).

DOI: 10.1039/D4QM00520A

This loss in through-plane thermal conductivity can also be attributed to the formation of voids between the printed layers (some are visible in **Fig. 5d<sub>3</sub>, d<sub>4</sub>**) which are responsible for creating macroscopic thermal contact resistances at the interface of two successive layers <sup>61</sup>. However, the orientation and alignment of AgNW@SiO<sub>2</sub>-20 in the 3D printed PC nanocomposites have a major influence on in-plane measurements. The 0/0 perpendicular cut (2) exhibited the best thermal conductivity, up to 3.48 ± 0.06 W m<sup>-1</sup> K<sup>-1</sup>, since all AgNW@SiO<sub>2</sub>-20 nanofillers were aligned in the testing direction. This corresponds to a fifteen fold increase in thermal conductivity with regard to neat PC. The alignment of the nanofillers in the printing direction is therefore responsible for very high anisotropy ratios (**Fig. 5e**). As expected, the 0/0 parallel cut (1) showed the lowest in-plane thermal conductivity (0.81 ± 0.01 W m<sup>-1</sup> K<sup>-1</sup>), since all the nanofillers were 90°-oriented regarding the in-plane testing direction. Interestingly, it was still 2.5 times higher than the through-plane thermal conductivity of the unidirectional sample (0.33 ± 0.03 W m<sup>-1</sup> K<sup>-1</sup>), even though the nanofillers were oriented in the same directions for both samples. This can be explained by a higher amount of interstitial voids between layers in the through-plane direction compared to the in-plane direction. Indeed, as the asymmetry of the printing beads is very high due to a higher extrusion width compared to the layer height (width/height ratio ≈ 4), there are in this case more layers (i.e. interfaces) in the through-plane direction than in the (x,y) plane. More interfaces means more voids, resulting in higher macroscopic interfacial thermal resistance. The 0/90 normal cut (3) in-plane thermal conductivity was found to be about the same as that of the bulk sample and intermediate between the in-plane values of the 0/0 parallel cut (1) and 0/0 perpendicular cut (2). This observation was expected since the 0/90 normal cut (**Fig. 5d<sub>3</sub>**) is a perfectly balanced mix between the 0/0 parallel cut (**Fig. 5d<sub>1</sub>**) and the 0/0 perpendicular cut (**Fig. 5d<sub>2</sub>**). This also demonstrates that the random orientation of the nanofillers in the bulk could be modeled by a balanced alternation of unidirectional layer orientations, since the thermal

Materials Chemistry Frontiers Accepted Manuscript

<figure><img src="image_25.png" /></figure>
conductivity measured in both cases is equivalent. As for the 0/90 45° cut, its in-plane thermal conductivity value was relatively close to that of the 0/90 normal cut, although slightly lower. The thermal conductivity values obtained in this work compare well to the results found in the literature for 3D printable, thermally conductive, and electrically insulating polymer nanocomposites (Table S3). To our knowledge, the original approach described herein delivers the best results to date (Fig. S15). Fig. 5f displays a proof of concept, constructed with both neat PC and PC-AgNW@SiO₂-20 and 3D printed with a rectilinear, unidirectional infill from base to top. Thermal imaging revealed thermal gradients for both composition and highlighted the superior heat dissipation and cooling abilities of the PC-AgNW@SiO₂-20 nanocomposite (Fig. S16). The orientation and alignment of 1D nanofillers can be a serious advantage if they are controlled and if the design of anisotropic parts is properly thought out. The properties of manufactured devices can thus be finely tuned and optimized by aligning the extrusion orientation with the desired preferential direction maximizing the heat dissipation properties.

## Conclusion

AgNW@SiO₂ core-shell nanostructures with continuous and conformal SiO₂ nanolayers of controlled thicknesses ranging from 5 to 43 nm have been successfully synthesized. Their use as thermally conductive and electrically insulating nanofillers to improve heat dissipation performances of a PC matrix was demonstrated, and the fabrication of 3D printable nanocomposite filaments at 3 vol% loading was achieved. This is, to the best of our knowledge, the first time that thermally conductive and electrically insulating polymer nanocomposites filled with core-shell 1D nanofillers are successfully 3D printed by FDM.

The silica nanolayer not only preserves the electrical insulation of the PC matrix thanks to the SiO₂ intrinsic properties, but also greatly improves the thermal conductivity of bulk PC nanocomposites (up to 2.08 ± 0.05 W m⁻¹ K⁻¹ at 3 vol% loading, a 190 % and 800 % improvement over PC-AgNW and neat PC, respectively). This dramatic thermal conductivity enhancement is

<figure><img src="image_26.png" /></figure>
attributed to a greatly improved filler-matrix cohesion and a much better distribution in the PC<sup>w Article Online</sup>
matrix.
DOI: 10.1039/D4QM00520A

The influence of the manufacturing process on thermal and electrical properties was also investigated. The most promising nanocomposite was 3D printed using the FDM technique. The resulting orientation and alignment of the nanofillers provided by the 3D printing technique led to remarkable in-plane thermal conductivities, up to 3.48 ± 0.06 W m⁻¹ K⁻¹ for the PC-AgNW@SiO₂-20 nanocomposite, while keeping the material electrically insulating (electrical resistivity > 10¹² Ω cm). This corresponds to a fifteen-fold increase in thermal conductivity with regard to the neat PC.

We thus demonstrate that the development of advanced core-shell nanowires can successfully lead to the 3D printing of high-performance and anisotropically controlled thermally conductive yet electrically insulating advanced materials. We believe this is a promising route to meet the ever-more challenging industrial needs for thermal management and on demand manufacturing of future heat dissipation devices.

# Experimental section

## Materials

Silver nanowires (AgNWs dispersion in isopropanol at 30 g kg⁻¹, diameter = 70 ± 5 nm, length = 10 ± 2 µm) were purchased from Protavic International (France). Tetraethyl orthosilicate (TEOS, 98.0 %), ammonium hydroxide solution (NH₄OH, 28.0-30.0 %) and chloroform (CHCl₃, 99.0 %) were purchased from Sigma-Aldrich (USA). Anhydrous EtOH was supplied by Carlo Erba (Germany). Polycarbonate (PC, grade Makrolon LED2245) was purchased from Covestro (Germany). Polytetrafluoroethylene (PTFE) membrane filters with a pore size of 1.2 µm were purchased from Sartorius (Germany).

## Synthesis of AgNW@SiO₂ core-shell nanostructures

Silica-coated silver nanowires (AgNW@SiO₂) were synthetized through a flexible sol-gel approach ⁶²⁻⁶⁴. The whole AgNW@SiO₂ synthesis procedure is detailed in the supplementary information file and summarized in the scheme displayed in Fig. 1a.

<figure><img src="image_27.png" /></figure>
## Preparation of PC-AgNW@SiO₂ nanocomposites

PC-AgNW@SiO₂ nanocomposites were synthetized by the solution mixing technique. Samples were prepared in order to obtain a 3 vol% loading rate of AgNW@SiO₂ in the polycarbonate matrix. Details about the preparation of the PC-AgNW@SiO₂ bulk nanocomposites can be found in the supplementary information file. The disk-shaped samples were obtained by hot pressing the flakes in a pressing die under an uniaxial hot press set at 270 °C (Carver 30T 4128, Carver Inc., USA). The nanocomposite synthesis process is summarized in the scheme presented in **Fig. 2a**. Note that PC-AgNWs bulk nanocomposites with pure AgNWs were prepared as control samples. For shaping PC-AgNW@SiO₂ nanocomposites in filament form, a similar process was employed as described in the supplementary information file and **Fig. S1**.

## 3D printing of neat PC and PC-AgNW@SiO₂ nanocomposite filaments

3D printing was carried out using an open source 3D printer (Hydra 16A, Hyrel 3D, USA). Details of the procedure can be found in the supplementary information file.

## Characterization

Characterization detailed methods are gathered in the supplementary information file. In particular, thermal conductivity values ($\lambda$) were calculated using the formula $\lambda = \alpha * \rho * C_p$, where $\alpha$ is the measured thermal diffusivity, $\rho$ is the measured bulk density and $C_p$ is the measured specific heat capacity.

## Acknowledgements

The authors acknowledge C. Brouard for helping in the silver titration experiments and analysis, A. Benayad for helping in the XPS characterization and M. Bey for helping in the breakdown voltage measurements. This work was supported by the Agence de l'Innovation de Defense (AID, France) through a PhD grant to Antoine Bodin.

<figure><img src="image_28.png" /></figure>

View Article Online
DOI: 10.1039/D4QM00520A
View Article Online
DOI: 10.1039/D4QM00520A

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DOI: 10.1039/D4QM00520A

The data supporting this article have been included as part of the Supplementary Information.

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