Ming Qi

Zou

Min Lei

Pei Xun Deng

Fan

Wei Li

Chen Fei Zhang

Sha Li

Lan Fan

Jean-François Jiang

Yong Feng Silvain

Lu

± Lei

Min Deng

Pei Xun Fan

Da Wei Li

Lan Jiang

Jean-Francois Silvain

Yong Feng Lu

Refractory vertically aligned carbon nanotube-boron nitride nanocomposites for scalable electrical anisotropic interconnects

Keywords: anisotropic, ceramic nanocomposites, vertically aligned carbon nanotube, electrical interconnects, boron nitride

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INTRODUCTION

Modern interconnect technology was introduced into the fabrication of printed circuit boards (PCBs) and very-large-scale integration (VLSI) at least two decades ago. [1][START_REF] Bohr | Interconnect Scaling-the Real Limiter to High Performance ULSI[END_REF][START_REF] Yee | Advantages and Challenges of 10-Gbps Transmission on High-Density Interconnect Boards[END_REF][4][5] It allowed PCBs to be designed with much higher routing density [START_REF] Romero | Advanced High Density Interconnection Substrate for Mobile Platform Application[END_REF] and revolutionized the VLSI industry by significantly increasing the performance of chips. [START_REF] Andricacos | Damascene Copper Electroplating for Chip Interconnections[END_REF] As feature sizes continuously shrink, interconnect technology demands much denser routing to decrease costs and avoid reliability risks. [START_REF] Romero | Advanced High Density Interconnection Substrate for Mobile Platform Application[END_REF] While copper (Cu) interconnects are dominating the market, they require barrier layers or special alloys to prevent poisoning when ramping up the routing density. 1 In addition, Cu interconnects suffer electromigration (EM) which worsens when the current density increases. [START_REF] Chai | High Electromigration-Resistant Copper/Carbon Nanotube Composite for Interconnect Application[END_REF] Carbon nanotubes (CNTs) have recently been recognized as a promising alternative to Cu interconnects. [START_REF] Wei | Reliability and Current Carrying Capacity of Carbon Nanotubes[END_REF][START_REF] Kim | Highly Aligned Scalable Platinum-Decorated Single-Wall Carbon Nanotube Interconnects[END_REF] It has been suggested that CNTs are the interconnect material of the future [START_REF] Kreupl | Carbon Nanotubes in Interconnect Applications[END_REF] based on their outstanding electrical properties as well as extraordinary failure current densities (>10 [START_REF] Wei | Reliability and Current Carrying Capacity of Carbon Nanotubes[END_REF] Acm -2 ). [START_REF] Wei | Reliability and Current Carrying Capacity of Carbon Nanotubes[END_REF][START_REF] Kim | Highly Aligned Scalable Platinum-Decorated Single-Wall Carbon Nanotube Interconnects[END_REF][START_REF] Nieuwoudt | Evaluating the Impact of Resistance in Carbon Nanotube Bundles for VLSI Interconnect Using Diameter-Dependent Modeling Techniques[END_REF][START_REF] Li | Bottom-up Approach for Carbon Nanotube Interconnects[END_REF][START_REF] Homma | Interconnection of Nanostructures Using Carbon Nanotubes[END_REF][START_REF] Naeemi | Monolayer Metallic Nanotube Interconnects: Promising Candidates for Short Local Interconnects[END_REF] In addition, CNTs have the ability to carry huge current density without heat sinks because of their exceptional thermal conductivity. [START_REF] Sahoo | Thermal Conductivity of Freestanding Single Wall Carbon Nanotube Sheet by Raman Spectroscopy[END_REF][START_REF] Hao | Brazed Carbon Nanotube Arrays: Decoupling Thermal Conductance and Mechanical Rigidity[END_REF] However, great difficulty exists in manufacturing perfectly aligned CNTs, which are beneficial for minimizing the conducting resistance. [START_REF] Behnam | Effects of Nanotube Alignment and Measurement Direction on Percolation Resistivity in Single-Walled Carbon Nanotube Films[END_REF][START_REF] Wu | Electrical Conduction and Percolation Behavior of Carbon Nanotubes/UPR Nanocomposites[END_REF] When stacking clustered CNTs, the intertube conduction becomes significant due to electron percolation, causing current leakage toward unwanted directions. [START_REF] Simoneau | Electron Percolation in Realistic Models of Carbon Nanotube Networks[END_REF] Consequently, CNT-dielectric composite systems have been introduced in order to manipulate the percolation by fine controlling the electrical percolation threshold (EPT) to prevent intertube conducting behavior, at the expense of high electrical conductivity, however, when compared to bare CNT networks. In such systems, consisting of a mix of dielectric and conductive materials, the formation of conductive pathways in the form of clusters of aligned CNTs inside the matrix is governed by classical percolation theories. [START_REF] Zeng | Characteristics of the Electrical Percolation in Carbon Nanotubes / Polymer Nanocomposites[END_REF][START_REF] Vigolo | An Experimental Approach to the Percolation of Sticky Nanotubes[END_REF] The conductivity of CNT-dielectric composite systems depends upon the alignment of nanotubes, with the highest conductivity being obtained via highly aligned CNTs. [START_REF] Behnam | Effects of Nanotube Alignment and Measurement Direction on Percolation Resistivity in Single-Walled Carbon Nanotube Films[END_REF] On this basis, a promising solution for addressing all of the stated issues is to combine vertically aligned carbon nanotubes (VACNTs) with appropriate dielectric materials.

As a thermally and chemically resistant refractory compound of boron (B) and nitrogen (N), boron nitride (BN) came to our attention. It exists in various crystalline forms that are isoelectronic to a similarly structured carbon lattice. [START_REF] Bu | Isoelectronic Doping of Graphdiyne with Boron and Nitrogen: Stable Configurations and Band Gap Modification[END_REF] Among those forms, hexagonal boron nitride (h-BN) demonstrates excellent anticorrosion ability, [START_REF] Zhang | Investigation of Hexagonal Boron Nitride as an Atomically Thin Corrosion Passivation Coating in Aqueous Solution[END_REF] thermal stability, [START_REF] Kostoglou | Thermal and Chemical Stability of Hexagonal Boron Nitride (h-BN) Nanoplatelets[END_REF] and, as a wide-bandgap semiconductor material, a decent dielectric property. [START_REF] Cassabois | Hexagonal Boron Nitride Is an Indirect Bandgap Semiconductor[END_REF] However, the fabrication of VACNT-BN remains a huge challenge. As we know, the average intertube distance in VACNT arrays is less than 10 nm, which greatly limits the infiltration of the majority of materials into the VACNT array. [START_REF] Li | Clean Double-Walled Carbon Nanotubes Synthesized by CVD[END_REF][START_REF] Zou | Thermally Stable and Electrically Conductive, Vertically Aligned Carbon Nanotube/Silicon Infiltrated Composite Structures for High-Temperature Electrodes[END_REF] Although thermal chemical vapor deposition (CVD), [START_REF] Zou | Thermally Stable and Electrically Conductive, Vertically Aligned Carbon Nanotube/Silicon Infiltrated Composite Structures for High-Temperature Electrodes[END_REF][START_REF] Jing | Coaxial Carbon@boron Nitride Nanotube Arrays with Enhanced Thermal Stability and Compressive Mechanical Properties[END_REF][START_REF] Jing | Thermal Conductivity Enhancement of Coaxial Carbon@Boron Nitride Nanotube Arrays[END_REF] and atomic layer deposition (ALD) [START_REF] Stano | Conformal Atomic Layer Deposition of Alumina on Millimeter Tall, Vertically-Aligned Carbon Nanotube Arrays[END_REF] have been reported to be effective in regard to the infiltration of materials into dense VACNT arrays, successful preparation of VACNT-BN nanocomposites with electrical anisotropy has rarely been realized.

In this work, we developed an efficient approach to the controllable fabrication of VACNT-BN nanocomposites, which utilized VACNTs as a conductive network and h-BN as an intertube shielding layer, via a single-step CVD process. Vertically aligned carbon nanotubes, rather than other carbon (C) nanomaterials, were chosen as the conducting framework due to their highly aligned structure which guarantees the conductive pathway after the infiltration of h-BN into the matrices and minimizes the thermal stress at elevated temperatures. [START_REF] Hao | Brazed Carbon Nanotube Arrays: Decoupling Thermal Conductance and Mechanical Rigidity[END_REF] The as-prepared VACNT-BN nanocomposites exhibited highly anisotropic conductivity in the network, the mechanism behind which was studied and attributed to the different electrical percolation thresholds in different directions as a nature of the VACNT-BN nanocomposites. The nanocomposites also have excellent high-temperature antioxidative ability (up to 1400 °C) and erosion resistance, suggesting that the prepared VACNT-BN nanocomposites have not only great potential for next-generation electrical interconnects but much more potential is yet to be explored.

METHODS

Fabrication of ultralong VACNTs via CVD:

Ultralong VACNTs were fabricated in a similar fashion as reported previously. [START_REF] Zou | Thermally Stable and Electrically Conductive, Vertically Aligned Carbon Nanotube/Silicon Infiltrated Composite Structures for High-Temperature Electrodes[END_REF] The growth of VACNTs was carried out in a two-zone CVD furnace (MTI Corporation, OTF-1200X) after appropriate catalysts were sputtered onto the substrate. Two different types of VACNT samples were prepared: VACNT films and cubicpatterned VACNTs.

Fabrication of VACNT-BN nanocomposites:

VACNT-BN nanocomposites were prepared via CVD. The VACNT samples were placed inside a tube furnace with customized gas inlets and were heated in an ammonia flow until 1100 °C was reached. Then, boron trifluoride was introduced into the chamber over a time period as long as 180 min. During the BN deposition, the chamber pressure was maintained at ~2.3 Torr.

Fabrication of VACNT-BN electrical conductivity test benches and I-V measurements:

VACNT-BN samples were thinned and lifted out of the bulk matrices via FIB (FEI, Helios NanoLab ™ 660). Electrode patterns were prepared via photolithography (SUSS MicroTec Group, MJB4 Mask Aligner) and Au was then sputtered by a magnetron sputtering system (AJA International, Inc., ATC-Orion 5 UHV) to obtain Au electrodes. The VACNT-BN lifted out were then transferred onto the Au electrodes. The I-V characteristics were measured by a probe station (Cascade MPS 150) and a semiconductor parameter analyzer (Agilent 4155C) at room temperature.

High-temperature antioxidative ability and thermal stability test: For antioxidative ability testing, the as-prepared h-BN film and VACNT-BN nanocomposites were oxidized in a muffle furnace (Thermolyne ™ Benchtop 1100 °C) at different temperatures until 1100 °C. Further oxidation from 1100 °C to 1400 °C was carried out in a tube furnace (MTI Corporation, GSL-1500X-RTP50), with an oxygen partial pressure of 100 mTorr. High-temperature gas erosion resistance of VACNT-BN was carried out by an oxy-acetylene torch, for as long as 126 min (Figure S16, supporting material).

Characterization: Raman spectral analysis of the samples was performed in a micro-Raman system (Renishaw, inVia™ Plus). A 514 nm Ar+ laser with a power of ~3 mW was used to excite the Raman scattering. Raman spectra and mapping were then collected through a 50× objective lens with an accumulated time of 10 s at each position. Scanning electron microscopy (FEI, Helios NanoLab 660) was used to study the morphologies of the as-prepared VACNT-BN nanocomposites. Energy-dispersive x-ray was carried out via this SEM system as well.

Transmission electron microscopy ( FEI, Tecnai Osiris S/TEM) was used to study the nanostructures of the grown BN and VACNT-BN samples. Energy-dispersive x-ray, EELS, and HAADF analysis were also carried out via this scanning TEM system.

RESULTS AND DISCUSSION

Design and Fabrication

Growth of thick boron nitride layers

In this study, thick boron nitride films were prepared via CVD. Hexagonal boron nitride film as thick as 11.8 μm (Figure S1, supporting material) was successfully obtained in a fashion similar to the one reported previously, [START_REF] Pierson | Boron Nitride Composites By Chemical Vapor Deposition[END_REF] with a growth time of 180 min, corresponding to a fast growth rate of ~ 4 μm/h. Interestingly, two different kinds of h-BN structures can be distinguished from the cross-sectional scanning electron microscopy (SEM) image shown in Figure S1b. To chemically examine the h-BN film grown, cross-sectional Raman mapping was obtained and is shown in Figure S2 to provide more structural and quality information on the as-grown, thick h-BN film.

The Raman mapping of peak intensity, peak position, and peak width of the as-grown h-BN were compared (Figure S2c-S2h, respectively). The peak intensity was strong in the top section (layered) of the h-BN and was slightly lower in the bottom section (particle-like) of h-BN. However, there was not much difference in the peak position and width of the E2g band across the cross-section, suggesting an overall uniform distribution of quality. [START_REF] Li | Atomically Thin Boron Nitride: Unique Properties and Applications[END_REF] In order to further confirm the crystal structure of these two different layers, a cross-sectional transmission electron microscopy (TEM) sample ~25 nm thick was prepared via focused ion beam (FIB), shown in Figure S3. As can be seen, the top BN layer shows a highly oriented, layered structure, as typical h-BN. On the contrary, the bottom BN layer contained many different polycrystalline orientations, which was confirmed to be amorphous BN. The electron energy loss spectroscopy (EELS) spectrum, shown in Figure S3e, further confirms the B-N structure.

Controlled growth of boron nitride on vertically aligned carbon nanotubes

We applied the same h-BN growth technique directly onto VACNTs instead of silicon dioxide/silicon (SiO2/Si). As previously reported, the infiltration of materials into dense VACNT arrays is extremely challenging; and the thermal CVD system was fine-tuned to ensure a homogeneous dispersion of h-BN inside the VACNT arrays. [START_REF] Zou | Thermally Stable and Electrically Conductive, Vertically Aligned Carbon Nanotube/Silicon Infiltrated Composite Structures for High-Temperature Electrodes[END_REF] Figure 1a illustrates the controlled fabrication process of VACNT-BN nanocomposites. In order to fabricate VACNT-BN samples, both VACNT films of different thicknesses (Figure 1a) and cuboid-patterned VACNTs (Figure 1b and 1c) were prepared for the h-BN infiltration. Optimized h-BN growth parameters were used for the infiltration; and the corresponding x-ray diffraction (XRD) results for VACNTs, h-BN film, and VACNT-BN (shown in Figure S4) were used to confirm the composite structure. As can be seen, both h-BN film and VACNT-BN exhibited a main feature of (002) while only VACNT-BN exhibits (004). It is reported that the (002) plane of h-BN exhibits similar properties of a graphenelike plane. [START_REF] Zhi | Large-Scale Fabrication of Boron Nitride Nanosheets and Their Utilization in Polymeric Composites with Improved Thermal and Mechanical Properties[END_REF] While the (004) plane of h-BN is parallel to the (002) plane, the XRD intensity of these two planes can be largely different due to different facets that are exposed in the crystals. [START_REF] Xue | Excellent Electrical Conductivity of the Exfoliated and Fluorinated Hexagonal Boron Nitride Nanosheets[END_REF] This indicates that the crystal alignment of h-BN film and that of VACNT-BN are different, which could be caused by a much more complicated surface condition of VACNTs for BN to grow.

The well-controlled growth of VACNT-BN yielded short (Figure S5) and ultralong (Figure S6 and S7) VACNT-BN nanocomposites. Figure S5 is a cross-sectional SEM image of an as-grown, short VACNT-BN sample. An excessive layer of h-BN can be seen above the lower VACNT-BN section. The VACNTs were heavily covered by h-BN as the fiber structures visually became much thicker. Figures S6 and S7 demonstrate the uniformity of the ultralong VACNT-BN nanocomposites along all directions. As an example, a 1.3-mm-long VACNT-BN cuboid was cut in half; and the cross-sectional SEM images for vertical (Figure S6) and horizontal (Figure S7) directions of the sample are presented. As shown in these results, the uniformity in macroscale turns out to be excellent.

Controlled intertube conduction manner

In our VACNT-BN system, the composite mixture was highly oriented by utilizing the VACNT as a template. Without insulating layers outside the VACNTs, current flowed along the direction of the VACNTs and intertube at the same time due to electron percolation (illustration Figure 1d). In order to achieve highly anisotropic conductivity, we engineered the VACNT-BN nanocomposites, based on the dielectric nature of h-BN. Dense VACNT arrays were used as the conducting network and building template, into which we infiltrated BN in a controlled manner. Hexagonal boron nitride served as an intertube shielding layer, increasing the electron percolation threshold of the VACNTs while maintaining the pristine conducting ability along the direction of the VACNTs.

Structural and Chemical Characterization

Although the graphene-like structure of the wall of VACNTs can act as a template for the nucleation of h-BN, achieving a stable and functional combination of these two materials remains a great challenge. [START_REF] Jing | Coaxial Carbon@boron Nitride Nanotube Arrays with Enhanced Thermal Stability and Compressive Mechanical Properties[END_REF][START_REF] Arenal | In-Situ Formation of Carbon Nanotubes Encapsulated within Boron Nitride Nanotubes via Electron Irradiation[END_REF][START_REF] Zhang | Carbon Nanotube and Boron Nitride Nanotube Hosted C60-V Nanopeapods[END_REF][START_REF] An | Linking Carbon and Boron-Nitride Nanotubes: Heterojunction Energetics and Band Gap Tuning[END_REF][START_REF] Walker | Growth of Carbon Nanotubes inside Boron Nitride Nanotubes by Coalescence of Fullerenes: Toward the World's Smallest Coaxial Cable[END_REF] To explore this matter, we first looked into the structure, in macroscale, of the VACNT-BN nanocomposites. Cross-sectional SEM images of them are presented in Figure 2. During the growth of h-BN, gas feedstocks first infiltrated into the VACNT array, forming h-BN layers which eventually accumulated on the top surface, resulting in the formation of a thick h-BN layer on top. As can be seen in Figure 2a and 2b, the h-BN layer on top exhibited a layered structure while the VACNT-BN region maintained its original vertically aligned feature (Figure 2d). The magnified SEM image (Figure 2b) gives a clearer view of the layered h-BN structure.

Different boundaries are outlined by the yellow dashed lines, indicating the evolution in size of h-BN layers, from small fractions to bigger plains. The small fractions are considered to be the consequence of direct growth on the top surface of the VACNT array, which was very rough. In addition, a clear boundary of VACNT-BN and h-BN regions can be identified (Figure 2c). To help better evaluate these different structures, Raman mapping of this cross-section area was carried out, as shown in Figure 2e-2g. A typical Raman spectrum of VACNT-BN is shown in the inset of Figure 2e. For the VACNT-BN region, the E2g band was separated from the combined D+E2g band via Lorentzian fitting. As can be seen from the E2g band intensity mapping, which represents the h-BN distribution along the cross-section, the top h-BN layer demonstrated a stronger signal than the VACNT-BN section. This was the result of peak shifting caused by the mixing of two materials. [START_REF] Miranda | Line Shape Analysis of the Raman Spectra from Pure and Mixed Biofuels Esters Compounds[END_REF] As for the G-band intensity mapping, which represents the VACNTs, it proved the existence of VACNTs under the heavy h-BN coating on the surface. A similar analysis was carried out on shorter VACNTs as well, and the results were consistent (Figure S8).

Although the above results visually and chemically confirmed that the h-BN was infiltrated inside the VACNT matrices, the combination between every single VACNT and the h-BN shielding layer was yet to be determined. Therefore, we looked further into this matter via various characterization techniques of scanning TEM (STEM). First, from a chemistry point of view, as can be seen in Figure 3b-3d, the results of energy dispersive spectroscopy (EDS) for (b) boron, (c) carbon, and (d) nitrogen mapping are presented. The high-angle annular dark field (HAADF) image of the sampled area is shown in Figure 3a as a reference. Strong signals for all three elements were obtained with good contrast. The distribution for both B and N was consistent with C, with no overlapping in the intertube space, indicating that the h-BN was only on the sidewall of the VACNTs. Taking full advantage of the better contrast provided by the HAADF image [START_REF] Jesson | Incoherent Im Aging of Crystals Using Therm Ally Scattered Electrons[END_REF][START_REF] Mehrtens | Measuring Composition in InGaN from HAADF-STEM Images and Studying the Temperature Dependence of Z-Contrast[END_REF] of the sample (Figure 3a), the h-BN coating on VACNT fibers was confirmed. To narrow down the scope and further look into one single VACNT, a HAADF (Figure 3e) image of one single VACNT is shown, alongside the EDS mapping of B (Figure 3f) and N (Figure 3g).as well. As can be seen, a single VACNT was suspended between surrounding structures, and both B and N elements were uniformly covering the VACNT wall.

However, some previous reports suggested that at elevated temperatures, C, B, and N can form complex chemical species that are not desired in our study. [START_REF] Wang | Porous Boron Carbon Nitride Nanosheets as Efficient Metal-Free Catalysts for the Oxygen Reduction Reaction in Both Alkaline and Acidic Solutions[END_REF][START_REF] Lei | Boron Carbon Nitride Nanostructures from Salt Melts: Tunable Water-Soluble Phosphors[END_REF][45][START_REF] Yap | B-C-N Nanotubes and Related Nanostructures[END_REF] To investigate this matter, EELS of the VACNT-BN sample of a larger area was collected. First, the low loss map of the nanocomposites is shown in Figure 3h, which suggests results consistent with the reported h-BN signature. [START_REF] Dobigeon | Ultramicroscopy Spectral Mixture Analysis of EELS Spectrum-Images[END_REF][START_REF] Cretu | Structure and Local Chemical Properties of Boron-Terminated Tetravacancies in Hexagonal Boron Nitride[END_REF] Another set of unique B-K signature peaks that only exist in BN was also identified, as shown in Figure 3i. The full EELS spectrum of the sample and additional features associated with C and N are shown in Figure S9. Interestingly, neither BN nor B-C-N features were identified in the spectrum. These results confidently indicate that the VACNT-BN nanocomposites we fabricated were free of undesired by-products and were only consist with VACNTs and h-BN. Furthermore, via close observation, white dots can be distinguished on the VACNT wall, the chemical nature of which, however, could not be determined solely by the HAADF image (Figure 3e). Hence, high-resolution transmission electron microscopy (HRTEM) of a single VACNT (Figure 4a) and VACNT-BN (Figure 4b) were obtained and compared with each other. Compared to a bare VACNT (Figure 4a) with clean walls, lots of additional layered features were identified on the wall of the VACNT-BN (Figure 4b) and the wall of the VACNT-BN is slightly distorted.

These features exhibited a similar graphene-like layered structure. To better evidence the existence of the h-BN, we took higher magnification HRTEM images of both VACNT-BN and h-BN film to compare with each other, and the figures have been added in the supporting material (figure S10) and is also shown below. High magnification HRTEM of the outside layers that are suspected to be h-BN is shown in figure S10a. As can be seen, in addition to the already identified layered structure, the atom-atom distance is roughly measured to be 2.6 Å and the distance consisting of two layers is around 5.9 Å. To verify these results, we compared them to the HRTEM of h-BN shown in figure S10b. It demonstrates a similar graphene-like layered structure, with a layer thickness of roughly 3~4 Å, which is consistent to both shown in figure 4b and reported results. [START_REF] Jing | Thermal Conductivity Enhancement of Coaxial Carbon@Boron Nitride Nanotube Arrays[END_REF][START_REF] Li | Atomically Thin Boron Nitride: Unique Properties and Applications[END_REF][START_REF] Nag | Graphene Analogues of BN: Novel Synthesis and Properties[END_REF] Combining this structural consistency and the elemental characterization provided by EELS and EDS, we are therefore positive that these graphene-liked layers outside the VACNT wall are h-BN. [START_REF] Arenal | In-Situ Formation of Carbon Nanotubes Encapsulated within Boron Nitride Nanotubes via Electron Irradiation[END_REF][START_REF] Walker | Growth of Carbon Nanotubes inside Boron Nitride Nanotubes by Coalescence of Fullerenes: Toward the World's Smallest Coaxial Cable[END_REF][START_REF] Yun | Selective-Area Heteroepitaxial Growth of h -BN Micropatterns on Graphene Layers[END_REF][START_REF] Yasuda | Graphitization Mechanism during the Carbon-Nanotube Formation Based on the in-Situ HRTEM Observation[END_REF] In addition, large h-BN nucleation sites were noted (dark round dots). The existence of these large spots is believed to correspond to the white features seen in the HAADF image (Figure 3e). Based on the location of these features, the graphene-like layers were determined to first form by the wall of a VACNT, which served as a template. These layers then served as transition layers for the larger scale nucleation of h-BN dots. This growth mechanism allowed a chemically and structurally stable combination of VACNTs and BN. [START_REF] Arenal | In-Situ Formation of Carbon Nanotubes Encapsulated within Boron Nitride Nanotubes via Electron Irradiation[END_REF][START_REF] Cretu | Structure and Local Chemical Properties of Boron-Terminated Tetravacancies in Hexagonal Boron Nitride[END_REF] However, as shown in Figure 4b, the wall of VACNT after BN growth has been slightly distorted due to the formation of these BN layers. Comparing to other reports with similar CNT-BN structures, [START_REF] Jing | Thermal Conductivity Enhancement of Coaxial Carbon@Boron Nitride Nanotube Arrays[END_REF][START_REF] Chang | Hexagonal Boron Nitride Coated Carbon Nanotubes: Interlayer Polarization Improved Field Emission[END_REF] our results show discontinued BN layers covering the VACNTs instead of a few smooth epitaxial layers. This discontinued formation of BN may have been caused by the purposed nucleation mechanism.

Electrical Anisotropy

Now that a kind of core-shell structure of the VACNT-BN nanocomposites prepared through our approach was confirmed, we explored its electrical potential performance towards interconnect applications. Two types of VACNT-BN test benches in similar sizes, with orthogonal orientations of VACNTs, were prepared. As can be seen in Figure 5a and 5b, the top view SEM images of the two test benches were (a) perpendicular to and (b) parallel to the VACNTs' direction, respectively.

The insets present the overall view of these two test benches. The test benches were welded onto gold (Au) electrodes in an all-around fashion to ensure good electrical connection. As can be seen in Figure 5a, the VACNTs were pointing downwards; and the current passed through the VACNTs horizontally. On the contrary, in Figure 5b, the VACNTs were pointing from left to right, parallel with the direction of the current. Interestingly, the current flowing through these two test benches exhibited a huge difference, of two orders of magnitude, in electrical conductivity.

Figure 5c compares the I-V curves of VACNT-BN measured on both test benches, where the blue and red lines represent the I-V parallel with the direction of VACNT and perpendicular to the direction of the VACNT, respectively. The electrical conductivity turned out to be 1060.43 Sm -1 along the VACNT direction, while only 4.43 Sm -1 measured perpendicular to the VACNT direction. Although the electrical conductivity of the VACNT-BN was significantly lower than bare multiwalled carbon nanotube (MWCNT) bundles, the electrical conductivity of which was reported to be ~2.810 [START_REF] Andricacos | Damascene Copper Electroplating for Chip Interconnections[END_REF] Sm -1 at a length of ~20 μm, [START_REF] Lekawa-Raus | Electrical Properties of Carbon Nanotube Based Fibers and Their Future Use in Electrical Wiring[END_REF] it showed superior electrical conductivity over CNT-polymer mixtures, whose conductivities were usually below 10 -3 Sm -1 , [START_REF] Zeng | Characteristics of the Electrical Percolation in Carbon Nanotubes / Polymer Nanocomposites[END_REF][START_REF] Nag | Graphene Analogues of BN: Novel Synthesis and Properties[END_REF] The sample also demonstrated good thermal stability and robust electrical performance after being heat cycled from 100 °C to 600 °C in argon flow, as shown in Figure 5d, which is significantly beneficial for practical applications. In addition, the red line, representing the current perpendicular to the direction of the VACNTs, showed some nonlinear characteristics. All of these results indicated the existence of different conductive pathways inside the VACNT-BN system. In an insulator-conductor system, multiple electron percolation thresholds could co-exist, due to the complex formation of the conducting network. [START_REF] Kovacs | Two Percolation Thresholds in Carbon Nanotube Epoxy Composites[END_REF] To better understand this phenomenon, the percolation effect of electrons in our VACNT-BN nanocomposites was analyzed. In our VACNT-BN system, the VACNTs were highly oriented and shielded by h-BN layers. Due to the dielectric nature of h-BN compared to the metallic VACNTs, the condition of electron transport was expected to be dominated by the pathways provided by the VACNTs themselves and mainly determined by compared various simulation works regarding conductive nanocomposites. [START_REF] Yasuda | Graphitization Mechanism during the Carbon-Nanotube Formation Based on the in-Situ HRTEM Observation[END_REF][START_REF] Chang | Hexagonal Boron Nitride Coated Carbon Nanotubes: Interlayer Polarization Improved Field Emission[END_REF] In general, with an increase in aspect ratio, an exponential decrease in the EPT was found, regardless of the waviness of the CNTs. [START_REF] Zeng | Characteristics of the Electrical Percolation in Carbon Nanotubes / Polymer Nanocomposites[END_REF][START_REF] Yasuda | Graphitization Mechanism during the Carbon-Nanotube Formation Based on the in-Situ HRTEM Observation[END_REF] As the typical diameter of our sample CNTs was ~ 20 nm (Figure 4b) and the length of the CNTs in the test benches (Figure 5a and5b) was ~20 μm, the aspect ratio was calculated to be 1000:1, which was significantly larger than the ones used in the simulation. [START_REF] Zeng | Characteristics of the Electrical Percolation in Carbon Nanotubes / Polymer Nanocomposites[END_REF][START_REF] Yasuda | Graphitization Mechanism during the Carbon-Nanotube Formation Based on the in-Situ HRTEM Observation[END_REF] Consequently, extremely conductive behavior along the direction of the VACNTs is expected.

As for the direction perpendicular to the VA CNTs, the introduction of BN on the wall of the VACNTs eliminated the contact between adjacent tubes. The extremely large band gap of BN made it a perfect insulating layer as well. In this case, the intertube conduction of electrons was extremely limited by the high barrier provided, which significantly increased the EPT, resulting in a significantly lower electrical conductivity along this direction. Meanwhile, the electron conduction along the VACNTs was not affected. Overall, VACNTs provided highly aligned electrical conductive pathways while the infiltrated h-BN in the nanocomposites induced the significant difference in EPT along different directions, [START_REF] Yun | Selective-Area Heteroepitaxial Growth of h -BN Micropatterns on Graphene Layers[END_REF] which comprehensively yielded the highly anisotropic electrical conducting performance of the prepared VACNT-BN nanocomposites.

As a contrast, we also measured the conductivity of bare VACNTs from both the parallel and perpendicular directions (Figure S11) and obtained very similar results although the geometrical aspect ratios were also significantly different for the two directions. It suggests that other factors, besides the geometrical feature of individual VA CNTs, take effect in determining the conductive behavior of bare VACNTs. In addition, the physical distortions of VACNTs can affect the electron transport in the CNT, [START_REF] Lekawa-Raus | Electrical Properties of Carbon Nanotube Based Fibers and Their Future Use in Electrical Wiring[END_REF] which helps explain the lower electrical conductivity of VACNT-BN (along CNT direction) than bare VACNTs. It was reported that for randomly aligned CNTs, measuring direction does not affect conductivity by too much. [START_REF] Simoneau | Electron Percolation in Realistic Models of Carbon Nanotube Networks[END_REF] For the bare VACNTs, multiple contact points actually existed among adjacent tubes. Thus, a complex and nondirectional conducting network formed, on which condition the conduction behavior of the electron was dominated not only by the pathways provided by individual VA CNTs but also by the intertube percolation. As a result, no obvious anisotropy in the electrical conducting behaviors was measured for the bare VACNTs.

It can be noticed that the level of electrical conductivity of VACNT-BN comparing with CNTs and other conductive material is substantially lower. In spite of that, when comparing to CNTmixture counterparts, such as CNT-Al2O3, [START_REF] Inam | Electrically Conductive Alumina-Carbon Nanocomposites Prepared by Spark Plasma Sintering[END_REF][START_REF] Flahaut | Carbon Nanotube-Metal-Oxide Nanocomposites: Microstructure, Electrical Conductivity and Mechanical Properties[END_REF] CNT-MgO 59 and CNT-Si3N4, [START_REF] Balázsi | Preparation and Characterization of Carbon Nanotube Reinforced Silicon Nitride Composites[END_REF] we still demonstrate competitive electrical conductivity. In our present research, we mainly intend to demonstrate an approach which can effectively infiltrate BN to VACNT forest to improve its thermal stability and electrical anisotropy. Although it is still early to demonstrate an application-ready VACNT-BN system due to the limited absolute electrical conductivity value currently reached, there are possible ways to improve the electrical conductivity in future works, such as: (1) reducing the amount of BN in the system. By controlling the amount of the BN infiltrated into the VACNT forest, the insulating layers between tubes will be thinner and some degree of conduction will be promoted while maintaining good enough anisotropic conducting ability; (2) reducing the growth rate of BN for finer control of BN layers formed outside of the VACNT walls, to avoid discontinuous BN formation and large defects. In addition, avoiding the formation of curvatures of VACNTs could effectively reduce strain in VACNTs, which could increase electron mobility in VACNTs 61 and (3) preparing better quality VACNTs. The quality improvement of bare VACNTs can enhance its electrical conductivity, which should also be able to enhance the electrical conductivity of the corresponding VACNT-BN nanocomposites.

Thermal Stability

Previous studies showed moderate thermal stability and reasonable antioxidative ability of VACNTs up until 600 °C. [START_REF] Zou | Thermally Stable and Electrically Conductive, Vertically Aligned Carbon Nanotube/Silicon Infiltrated Composite Structures for High-Temperature Electrodes[END_REF][START_REF] Frank | Oxidation Stability of Multiwalled Carbon Nanotubes for Catalytic Applications[END_REF] Meanwhile, h-BN demonstrated very good antioxidative ability at elevated temperatures. [START_REF] Kostoglou | Thermal and Chemical Stability of Hexagonal Boron Nitride (h-BN) Nanoplatelets[END_REF] To verify this, the antioxidative stability of our as-grown h-BN film was tested in air. Raman spectra were collected after each oxidation process. As shown in Figure S12a, samples exhibited no evident change as the oxidation temperature increased to even 1100 °C. We further investigated the oxidation temperature dependence of peak position and full width at half maximum (FWHM) of the E2g peak of h-BN. No obvious change was observed in either FWHM or the peak position, indicating the high quality and high-temperature oxidation resistance of the h-BN films grown by our approach.

The effect of oxidation time on the high-temperature oxidation resistance of the as-grown h-BN films was also studied by treating them at 1100 °C in the air for an extended period (Figure S13). The thickness of the h-BN films was evaluated after each oxidation process via crosssectional SEM. The original thickness of the h-BN film was ~ 11.8 μm (Figure S13a). After being oxidized at 1100 °C for 5 min, the film thickness decreased slightly to ~ 11.5 μm (Figure S13b).

When the oxidization time was extended to 40 min, the film thickness was decreased to ~ 8 μm, suggesting a slow oxidation rate at 1100 °C. In spite of this, the oxidation only happened on the upper section of the layered h-BN, indicating it was a good candidate for the protective layer for the VACNTs.

Although the antioxidative ability of h-BN has been determined, the effect of infiltrating h-BN into VACNT arrays is still unknown. Therefore, we further investigated the antioxidative ability of as-prepared VACNT-BN nanocomposites. The VACNT-BN nanocomposites were first oxidized at 1100 °C for a long treatment time of 40 min in air. Figure S14a Encouragingly, the oxidation process demonstrated behavior similar to the h-BN layer, which showed that the decrease in thickness only happened in the upper section of the thick h-BN film.

Meanwhile, the lower VACNT-BN area was not affected at all. An enlarged SEM image in Figure S14c shows the layered h-BN film near the top surface to be covered with multiple visible rough oxidation sites, suggesting that the oxidation started from the top surface of the h-BN. Scanning electron microscopy images in Figure S14d-S14f show no visual evidence of oxidation in the original VACNT area, indicating a good protection effect provided by the infiltrated h-BN. The thickness decrease under the high-temperature, oxidation condition in the upper h-BN film was expected and could be improved by increasing the thickness of the upper h-BN further, if needed.

It's worthy of noting that, during the measurement of electrical conductivity of our VACNT-BN nanocomposites, we need to expose the CNT tips to ensure good electrical contact. However, during practical applications, BN will be deposited to cover not only the VACNTs, but also the electrical contact and corresponding mating parts, so that its refractory ability and protective effect can be sufficiently utilized. Since h-BN can be stable up to much higher temperatures, the antioxidative ability of the VACNT-BN nanocomposites was examined under even harsher conditions, including ultrahightemperature oxidation treatment (up to 1400 °C) and an Ultrahigh-temperature gas erosion test. To better visually observe the oxidation results, VACNT-BN samples in the cubic pattern, instead of films, were used and oxidized from 1100 °C to 1400 °C, for 5 min at each temperature, with oxygen partial pressure of 100 mTorr (Figure 6a-6d, respectively). Magnified SEM images of the top surface are shown in Figure 6e-6h, corresponding to Figure 6a-6d. As can be seen, the originally smooth surface became rough as the oxidation temperature increased, and significant surface morphology change was observed beginning at 1300 °C (Figure 6g). Meanwhile, the height of each VACNT-BN cuboids started to drastically reduce after 1300 °C (Figure 6i), indicating the loss of VACNTs during the oxidation. It has been reported that h-BN are prone to oxidize under this condition accompanied by a slight weight gain and a decrease in size. [START_REF] Jacobson | High-Temperature Oxidation of Boron Nitride[END_REF] The shrinkage of the outside BN caused thinning of the protective BN layer, eventually exposing the inside VACNTs to the environment. In addition, Raman characterization of samples in figures 6a-6d are shown in figure S15. Comparing to the Raman spectrum of the sample before any oxidation treatment (figure S15a), the combined feature D+E2g around 1370 cm -1 demonstrate no significant loss in intensity and a stable peak position. The G-band associating to CNTs, which is around 1580 cm -1 , gained slightly change in intensity and became more pronounced after oxidation. The peak position of Gband maintained relatively stable after 5 min of oxidation at each temperature. When prolonging the oxidation duration to 30 min at 1400 °C, the G-band blue-shifted to 1601 cm -1 . One of the possible reasons for such a blueshift of the G-band is that the thermal expansion coefficient of CNT (positive) and BN (negative) are mismatched. When heating for a longer time, the BN layers compress the inside CNTs, causing the increase in the Raman shift. Despite this, the overall chemical composition stayed unchanged, which indicates that the outside BN layers can protect the inside VACNTs well at temperature as high as 1400 °C. The result of oxidation of VACNT-BN nanocomposites is a material loss rather than chemical change since carbon and boron does not react to form any side products like B-C-N. Thus, the oxidation cannot be reflected in the Raman spectroscopy. It is this material loss caused by oxidation as well as the shrinkage of BN under oxidative environment who should account for the significant decrease of the height of VACNT-BN nanocomposites after 1300 °C.

An ultrahigh-temperature gas erosion test was performed as well. Figure S16a S16d. As can be seen, the main features, namely, the D-and G-band for VACNT, as well as the E2g band for h-BN, presented no apparent changes, confirming good erosion resistance of the prepared VACNT-BN nanocomposites against oxygen contained in the ultrahigh temperature, high-velocity gas flow. As a comparison, no on par stability performance for the CNT-polymer-based nanocomposites has been reported as yet. The excellent stability of VACNT-BN as electrical interconnects demonstrates great potential for high-power and high-current density applications, such as insulated-gate bipolar transistors (IGBT).

CONCLUSIONS

In summary, a novel approach has been demonstrated for the finely controlled fabrication of refractory VACNT-BN conductive nanocomposites. Ultralong VACNTs (up to 3 mm), as superaligned conducting pathways and building framework templates, were prepared based on which uniformly infiltrated VACNT-BN composite structures were successfully obtained via the one-step CVD process. The as-fabricated VACNT-BN nanocomposites exhibited significant anisotropic electrical conducting performance, excellent thermal stability, and the extraordinary ability to withstand various extreme conditions (e.g., ultrahigh-temperature oxidation, and highvelocity hot gas erosion). The h-BN layers which were uniformly infiltrated into the VACNT arrays significantly reduced the intertube electron percolation, providing excellent stability enhancement at the same time. These attributes of VACNT-BN nanocomposites demonstrated great potential to be superior to many CNT-based counterparts for high density interconnect applications. In addition, the CVD process proposed in this work requires minimal complexity and effort in sample preparation, which is promising for the development of a cost-effective, turnkey solution to fabrication of scalable electrical interconnects. Our future work will explore the potential applications for VACNT-BN nanocomposites and their properties under extreme conditions.

Figure 1 .

 1 Figure 1. Illustration of controlled fabrication process and electron percolation mechanism. (a) The controlled fabrication process of VACNT-BN. Two cross-sectional SEM images demonstrate two different types of VACNT-BN film samples. (b) Optical image and (c) SEM image of cuboid-patterned VACNT-BN sample. Inset in (b) shows one single magnified VACNT-BN cuboid. (d) Different electron percolation conditions in VACNT and VACNT-BN. Scale bars for (a) 20 μm, (b) 1 cm, and (c) 50 μm.

10

 10

Figure 2 .

 2 Figure 2. Microscale structural and chemical analysis of prepared VACNT-BN nanocomposites. (a) Low magnification cross-sectional SEM images of VACNT-BN. Scale bar: 10 μm. (b-d) High magnification crosssectional SEM images corresponding to the top, interfacial, and bottom area of (a), respectively. Scale bar: 5 μm. (e) Optical microscopic image of the sample. (f,g) Cross-sectional Raman intensity mapping of (f) E2g band and (g) G band of the area shown in (e). Scale bars for (e-g): 10 μm.

Figure 3 .

 3 Figure 3. Nanoscale elemental analysis of prepared VACNT-BN nanocomposites. (a) HAADF image, (b) boron, (c) carbon, and (d) nitrogen EDS mapping of VACNT-BN fibers. Scale bars: (a-d) 1 μm. (e) HAADF image, (f) boron, and (g) nitrogen EDS mapping of a single VACNT-BN. Scale bars: (e-g) 40 nm. (h) EELS low loss spectrum of VACNT-BN. (i) BN signature band in EELS spectrum of the VACNT-BN sample.

Figure 4 .

 4 Figure 4. Nanoscale structural analysis of prepared VACNT-BN nanocomposites. HRTEM images of single (a) bare VACNT and (b) VACNT-BN nanocomposites. Scale bars: 5 nm.

Figure 5 .

 5 Figure 5. Electrical anisotropy of prepared VACNT-BN nanocomposites. (a,b) VACNT-BN test benches for I-V analysis where the direction of VACNTs is (a) perpendicular to and (b) parallel with the direction of the current flow. Insets present low magnification SEM images of the test benches. (c) I-V characteristics in log scale of (a) red and (b) blue. (d) I-V characteristics of (b) after annealing in argon flow from 100 °C to 600 °C. Scale bars for (a) and (b): 500 nm. Scale bars for insets of (a,b) : 15 um.

  the geometrical feature of individual VA CNTs. Natsuki et al. investigated different electrical percolation behaviors as a result of different orientations and aspect ratios of fillers. Lu et al.

  and S14b compare the change in sample thickness (a) before and (b) after the high-temperature oxidation treatment. The thickness of the sample decreased by ~ 7 μm in total after 40 min of oxidation at 1100 °C.

Figure 6 .

 6 Figure 6. Antioxidative ability of prepared VACNT-BN nanocomposites. (a-d) SEM images of VACNT-BN cuboid samples after oxidation at (a) 1100 °C, (b) 1200 °C, (c) 1300 °C and (d) 1400 °C. (e-h) Magnified SEM images of the top area of the samples from (a) to (d), respectively. (i) Height change of VACNT-BN cuboid after oxidation. (j) Raman spectrum of VACNT-BN after being oxidized for 30 min at 1400 °C. Fitted lines represent individual peaks of interests.

  Figure S16c inset. Raman spectra at slightly different stages of the erosion test are shown in Figure

ACKNOWLEDGMENTS

This research was financially supported by U.S. Department of Energy [NETL: DE-FE0023061], and the Nebraska Center for Energy Science Research. Manufacturing and characterization analyses were performed at the NanoEngineering Research Core Facility (part of the Nebraska Nanoscale Facility), which is partially funded by the Nebraska Research Initiative.

ASSOCIATED CONTENT

Supporting Information. The Supporting Information is available free of charge on the ACS Publications website at http://pubs.acs.org.

AUTHOR INFORMATION * Corresponding Author

E-mail: ylu2@unl.edu E-mail: fanpeixun@gmail.com

CONFLICTS OF INTEREST

The authors declare no conflict of interest.

NOTES

± Authors contributed equally.

Refractory Vertically Aligned Carbon


Adrien Morvan

Jean-Luc Grosseau-Poussard

Nathalie Caillault

Florence Delange

Sophie Roure

Pascal Lepretre

Jean-François Silvain
email: jean-francois.silvain@icmcb.cnrs.fr

Powder processing methodology for fabrication of Copper/Graphite composites with enhanced thermal properties

Keywords: A -Metal-matrix composites (MMCs), B -Thermal properties, E -Powder processing methodology

This study focuses on the powder processing methodology (PPM) for the fabrication of metal matrix composite using Graphite flakes (Grf) reinforced Copper (Cu). The manufacturing route involved 1) a reductive treatment of Grf powder to purify and increase its quality, 2) the mixing of the Cu and the Grf (0 to 40 vol.%) powders with a fast and efficient resonant acoustic (RA) mixer, and finally 3) the cold and hot-pressing of composite powders. Comparison of PPM is made with literature and a usual method used in our laboratory. The quality of Grf after different steps was analyzed by Raman spectroscopy and phase compositions and microstructure of the composite materials were analyzed by XRD and SEM, respectively. It was shown that this new PPM demonstrated better composite structural and thermal properties even at low volume fraction of Grf with a maximum of 630 W.m -1 .K -1 obtained with 40 vol.% of Grf at 70 °C. This was possible thanks to our new PPM which allows getting lower defects in the Carbon lattice of Grf and oxygen content in the final composite materials.

Introduction

Recently, the adaptive physical properties of Carbon/Metal Matrix Composites (MMCs) is receiving increasing attention. Addition of Carbon (C) such as Graphite flakes (Grf) to metallic matrix such as Copper (Cu) reduces its density and increases its thermal conductivity (TC) measured parallelly to the basal plane of Grf.

The fabrication of such composite materials is usually carried out by powders metallurgy process in two main steps: powders mixing and consolidation via sintering technique. In most cases, homogenization of metal and reinforcement powders was obtained after mixing by mechanical [START_REF] Firkowska | The Origin of High Thermal Conductivity and Ultralow Thermal Expansion in Copper-Graphite Composites[END_REF] [2] [START_REF] Wang | Effect of the content of ball-milled expanded graphite on the bending and tribological properties of copper-graphite composites[END_REF] or wet mixing [START_REF] Kováčik | Thermal conductivity of Cu-graphite composites[END_REF]. These processes are long or/and induce degradation of reinforcements by chemical contamination, plastic deformation or by local overheating.

Analyses performed by Raman spectroscopy on Grf, obtained after ball-milling, have demonstrated that the structure of Grf was damaged as the ID/IG ratio increased after mixing step and densification [START_REF] Firkowska | The Origin of High Thermal Conductivity and Ultralow Thermal Expansion in Copper-Graphite Composites[END_REF].

In this work, a new powder processing methodology (PPM) was investigated to avoid creation of defects in Grf and therefore to improve the thermal properties of Grf and consequently the thermal properties of the Cu/Grf composite materials. We have investigated how 1) a reductive treatment of the Grf, prior to mixing step, tend to reduce the oxygenated sp 3 -C content and 2) how a new and fastest mixing method homogeneously disperse the Grf inside the Cu matrix without generated any degradations of the mixed powders. This mixing method, which is based on resonant acoustic (RA) mixing, is already used by pharmaceutical industries to mix powder bled. This non-contact mixing technology helps to improve quality at the edges of Grf thanks to a multitude of micro-mixing cells throughout the vessel. The effect of this new PPM on the TC of Cu/Grf composite materials is demonstrated and discussed. Comparison of composites TC obtained with conventional 3-dimensional (3D) mixing process has been done. Using this new PPM, TC of the Cu/40 vol.% Grf composite material, after consolidation using conventional hot-pressing method, showed in-plane TC up to 630 W.m -1 .K -1 at 70 °C.

Experimental sections: a) Materials

Cu particles with dendritic shapes (Figure 1) have a d50 close to 30 µm and are supplied by ECKA ® Granules (Germany, GmbH). The dendritic powder has been chosen for this study because of its high purity and its dendritic shape which lead to easy grain deformability and therefore high final density. Figure 2 shows SEM micrographs of raw natural Grf with particles mean lateral size close to 500 µm and an average thickness of 30 µm. They are provided by Yanxin-Graphite Co., Ltd. As it was already discussed by Ben Liu et al. [START_REF] Liu | Effect of graphite flakes particle sizes on the microstructure and properties of graphite flakes/copper composites[END_REF], Grf particles size could affect the final physical properties (such as TC) of the composite material. Thereby, prior to be used, Grf were sieved to obtain a homogeneous powder lateral size, ranging between 355 and 400 µm.

b) Reduction of graphite powder

To perform the heat treatment under reductive atmosphere (Ar/H2(5%) with a flow of 200 mL/min), a certain amount of Grf was placed in an alumina pod and then placed in the center of a tubular furnace. The furnace was heated at 6 °C/min to the bearing temperature (200 °C, 400 °C and 800 °C) and then kept for an hour at annealing temperature. The decrease of temperature was obtained by natural cooling of the furnace.

c) Powder mixing

Composite powders (Cu/Grf) were mixed in a classical 3D mixer (Turbula ® Shaker/Mixer Model T2C, Germany) for 120 min at 40 RPM. Another mixing process was evaluated and compared to the 3D mixing, the RAM (Labram, Resodyn, USA). This device was operating at 60 Hz, during 30 seconds with an acceleration of 100 g.

d) Composite materials

Three sets of composite materials were sintered. Set A, Cu/Grf after 3D mixing; Set B after RA mixing and Set C, Cu/Grf (Ar/H2(5%)) fabricated after RA mixing with an additional step in the process: reductive treatment of Grf under Ar/H2(5%) before mixing. The mixture, with a volume fraction of Grf ranging between 0 to 40 %, was cold pressed in a carbon mold of 20 mm in diameter and then placed in a Thermolab ® press. All Sets of powders were heated at 650 °C using induction system, under a uniaxial compressive stress of 60 MPa for 45 min. The temperature was monitored via a K-type thermocouple located close to the sample.

e) Microstructural and chemical characterizations

The FT-IR spectra was recorded by grinding spectroscopy grade KBr with Grf using a BRUKER Equinox 55 FT-IR spectrometer in the range of 400 to 4000 cm -1 at 4 cm -1 resolution. Thermogravimetric analysis (TGA) was carried out on Grf using a TA Q600 thermal analyzer (SDT Q600, TA Instrument Co., USA) under a slightly hydrogenated atmosphere (Ar/H2(5%)), over a temperature range of 30 °C to 1000 °C with a heating rate of 6 °C.min -1 .

A ThermoFisher Scientific K-ALPHA spectrometer was used for surface analysis of Grf with a monochromatized Al Kα source (h = 1486.6eV) and a 200 microns spot size. A pressure of 10 -7 Pa was reached in the chamber when transferring the samples pressed onto indium foil.

The full spectra (0 -1100 eV) were obtained with constant pass energy of 200 eV and high-resolution spectra at a constant pass energy of 40 eV. Charge neutralization was applied during analysis. C1s, O1s high resolution spectra were fitted and quantified using the AVANTAGE software provided by ThermoFisher Scientific (Scofield sensitivity factors used for quantification).

Grf powder X-ray diffraction (XRD) patterns were collected on a PANalitycal X'pert PRO MPD diffractometer in Bragg-Brentano θ-θ geometry equipped with a secondary monochromator and X'Celerator multi-strip detector. Each measurement was made with an angular of 2θ varying between 8 and 80 with a 2θ step size of 0.017 °. The Cu-Kα radiation was generated at 45 KeV and 40 mA (λ = 0.15418 nm). The patterns were analyzed using EVA software (Bruker).

Grf was characterized after each steps of the composite fabrication process by Raman spectroscopy. The Raman analyses were done with a high-resolution Horiba spectrometer equipped with a CCD detector; its precision is equal to 0.1 cm -1 . Measurements were taken in a backscattering micro-configuration. The wavelength of the source is at λ = 532 nm, whose maximum power is equal to 10 mW. The laser power is chosen between 5 and 10 mW to avoid local heating of the samples. This power range is sufficient to obtain optimum resolution. All our experiments were performed at room temperature (RT) with a 50x microscope objective (leading to an analyzed surface area close to 1.6 μm 2 .) Positions of the selected bands have been calculated using Pseudo-Voigt and Lorentzian single and multipeak fitting (Fityk 0.9.8 curve fitting and data analysis program) after background subtraction.

Prior to SEM analysis and thermal measurements, due to anisotropic microstructure and properties, samples were machined to obtain two pellets (Ø 6 × 3.5 mm) revealing transverse (out-of-plane) and in-plane directions of Grf (Figure 3). The starting powders and bulk composites (orientation distribution) microstructures was analyzed by Scanning Electron Microscopy (SEM, VEGA II SBH, TESCAN) equipment. To reveal the microstructure, Cu/Grf composites were prepared using mechanical polishing with waterproof abrasive silicon carbide papers and polishing cloth.

Considering that the composite materials only contain closed porosity (ranging from 0.2 and 2 %), the relative density of the Cu/Grf composites was measured using the Archimedes principle.

Amount of oxygen was measured for pure sintered Cu and composites Cu/Grf. Hot extraction analyses were carried out to determine the elements concentration of materials, using EMGA-621W oxygen/nitrogen/hydrogen analyzer from Horiba. The estimated values were quantified, after combustion in a chamber furnace, by a thermal conductivity detector (TCD) and a nondispersive infrared sensor (NDIR).

f) Thermal characterization

The thermal diffusivity α of composite materials was measured parallel and perpendicular to the compressive axis with a LFA 457 Nanoflash laser thermal conductivity apparatus (NETZSCH, Germany). The specimen dimensions were 6 mm x 3.5 mm (diameter x thickness). A graphite coating layer was applied to the pellet's materials before thermal diffusivity measurements. The measurements were performed from RT to 130 °C.

The specific heat Cp of the composite was calculating using equation 1:

Cp = fCu.Cp Cu + fGrf.Cp Grf , (1)
Where fCu, fGrf and Cp Cu , Cp Grf are the weight fractions and the specific heat capacities of Cu and Grf, respectively. Heat capacity was measured by Differential Scanning Calorimetry from RT to 300 °C (DSC 8000 Pyris Diamond PerkinElmer). Grf and Cu powders have a heat capacity at 70 °C of 836 J.kg -1 .K -1 and 393 J.kg -1 .K -1 , respectively.

Thermal conductivity values were obtained by using equation 2:

K = ρc.Cp.α, ( 2
)
where K is the thermal conductivity (W.m -1.K -1 ) of the composite, α is the thermal diffusivity (mm².s -1 ), ρc (g.cm -3 ) is the bulk density and Cp (J.kg -1 .K -1 ) is the specific heat capacity of the composite. A rule of mixture is applied to determine theoretical density ρ of Cu/Grf composites. presented in Table 1. It was obtained by calculating the integrated area of the deconvoluted components in the C1s spectra of Figure 6.a. This material is slightly hydrophilic as it was already found by FT-IR and is prone to react with his environment. A very small amount of Chlorine (Cl) and Sulfur (S) has been detected. The amount of contamination could arise from a surface adhesive as it was mentioned by Ben Liu [START_REF] Liu | Effect of graphite flakes particle sizes on the microstructure and properties of graphite flakes/copper composites[END_REF].

Results and discussion

Both Raman spectra on Figure 7.a)1 and Figure 7.a)2 from center and edges, respectively,

show the four most prominent features in Raman spectra of graphitic materials [START_REF] Pimenta | Studying disorder in graphite-based systems by Raman spectroscopy[END_REF]. The G band appears around 1580 cm -1 and corresponds to sp²-C network. The D band around 1350 cm -1 and D' band, around 1620 cm -1 , appear when sp 3 -C and sp-C are introduced in the C-lattice.

Finally, the 2D band, which is the overtone of D band, appears around 2720 cm -1 . From Figure 7.a), one can see that more defects are introduced in the C-lattice at the flake edges as bands D and D' are more intense in the spectra 2 than in the spectra 1 (related to the flake center).

b) Effect of reductive treatment on Graphite flakes reinforcement

To remove most of the oxygen sp 3 -C species from Grf, a reductive treatment under Ar/H2(5%)

was chosen instead of a chemical treatment [START_REF] Liu | Effect of graphite flakes particle sizes on the microstructure and properties of graphite flakes/copper composites[END_REF]. To determine the desorption kinetics under reductive atmosphere, TGA of Grf up to 1000 °C have been performed. Figure 8 displays the weight loss in three discrete stages. In accordance with the analysis obtained by [START_REF] Sun | Reduced Graphene Oxide Reinforced 7075 Al Matrix Composites: Powder Synthesis and Mechanical Properties[END_REF] [8], the weight loss (close to 0.01 %) in the first stage before 150 °C is mainly caused by the elimination of physiosorbed and interlamellar water molecules. In the second stage between 150 °C and 350 °C, significant mass loss (close to 0.59 %) is primarily attributed to the thermal decomposition of covalently bonded oxygen as by-products of the reduction process. Finally, in the last stage the mass loss (close to 0.6 %) can be attributed to the removal of more stable Oxygen-containing functional groups and unstable Carbons remaining in the structure. Total amount of Oxygen species desorbed during annealing treatment under Ar/H2(5%) is close to 1.2 wt.%. This amount is not negligible as it is known that presence of Oxygen in the final composite could lead to a decrease in TC [START_REF] Jakani | Effet des impuretés sur les mécanismes de recristallisation du cuivre tréfilé[END_REF]. Therefore, it is necessary to know which bearing temperatures could allow getting lower amount of oxygenated sp 3 -C without degrading the Clattice of Grf.

The (002) d-spacing for Grf after several reductive treatments range from 3.377 Å (as-received) to 3.360 Å which compares to 3.356 Å for single crystal graphite (Table 2). The slightly larger d-spacing confirms the presence of sp 3 -C species as found, for example, by XPS analysis.

Adams [START_REF] Adams | Characterization of high thermal conductivity carbon fibers and a self-reinforced graphite panel[END_REF] reported that TC of Carbon fibers and self-reinforced graphite panel is more correlated with the graphite inter-basal-plane spacing d(002) than with the crystallite size (Lc or La) calculated from the Scherrer equation. Therefore, small d(002) leads to higher TC of Carbon species. In addition, improved layer plane stacking is consistent with the removal of defects such as vacancies, interstitials and impurities. Thus, the decrease of the d(hkl) of the (002) plane with the thermal treatment can be associated with the increase of the TC of the Grf. According to this correlation between the TC and d(002) spacing found by Adams, the estimated TC of the Grf after reductive treatment at 400 °C is close to 1100 W.m -1 .K -1 , comparing to about 500 W.m -1 .K -1 for as-received Grf (Erreur ! Source du renvoi introuvable.).

The integrated intensity Raman spectroscopy AD/AG ratio is widely used for characterizing defects in graphitic materials [1][6]. By comparing starting and annealed materials, it is possible to quantify defects introduced to Grf after each step of the process. Results, summarized in Erreur ! Source du renvoi introuvable. (average of at least six acquisitions), show that the reduction process helps to decrease the AD/AG ratio especially on Grf edges. The integrated intensity ratio AD/AG for the as-received Grf is 0.13 on the center and increase to 0.36 on the edges (Table 3). This difference of defects is mostly due to broken and folded edges of the flakes as it can be seen on Figure 11 and was already observed by M.A Pimenta on nanographite platelets [START_REF] Pimenta | Studying disorder in graphite-based systems by Raman spectroscopy[END_REF]. Nevertheless, this ratio shows that the as-received Grf possesses quite low defects and should display good thermal properties. With annealing treatment, the decrease of defects amount confirms the restoration of π-π conjugation through the removal of oxidative debris of Grf. At 200 °C, desorption of water and first carbonyls are taking place but are not yet complete while at 800 °C the graphitic network is starting to deteriorate du to destruction of C-C bonds. These two assumptions were confirmed by the increase of AD/AG ratio and the decrease of the WG and W2D positions shift on both localizations. However, the minimum number of defects of the Grf is obtained after reduction treatment at 400 °C. This is coherent with XRD measurements where the treatment at 400 °C was optimal in term of dspacing and theoretical TC. Therefore, this bearing temperature was chosen for the thermal treatment of Grf as it allows getting the lower defects amount in the C-lattice and should induce higher TC of the final composite materials. [START_REF] Mellmann | The transverse motion of solids in rotating cylinders-forms of motion and transition behavior[END_REF]. For the 3D mixing used in this work, a Fr equal to 0.9 has been calculated for a filling degree of the vessel of 0.5. This number permits to say that the motion, which is taking place in the mixing vessel, is a cascading type [START_REF] Mellmann | The transverse motion of solids in rotating cylinders-forms of motion and transition behavior[END_REF]. This flow mode gives rise to high shear mixing between powders as well as between powders and vessel walls.

Therefore, frictional heat is created during mixing and increase with mixing time. This overheating, in combination with the 2 hours mixing time, tends to deteriorate the quality of Grf. In contrast, RA mixing time is much smaller (less than a minute). In addition, the use of low frequency (60 Hz) induces high intensity energy mixing cells which allow a random dispersion of powders throughout the mixing vessel and to reduce frictional heat and processing times [START_REF] Tanaka | Verification of the mixing processes of the active pharmaceutical ingredient, excipient and lubricant in a pharmaceutical formulation using a resonant acoustic mixing technology[END_REF]. Table 4 shows that both peak positions of G and 2D, in cm -1 , taken on the edge of Grf, decrease after RA mixing. A. Ferrari et al, have concluded from Raman spectroscopy analyses on graphene/graphite, that a negative shift (in cm -1 ) of both G and 2D peak positions correspond to the introduction of tensile strain while a positive shift corresponds to a compressive strain in the C-lattice [START_REF] Ferrari | Raman spectroscopy of graphene and graphite: Disorder, electron-phonon coupling, doping and nonadiabatic effects[END_REF]. From that results, it could be considered that a larger amount of tensile strain has been introduced inside the edges of Grf during the RA mixing which could help to decrease amount of broken parts and folded flakes. After uniaxial hot-pressing, the AD/AG ratio increases from of 0.12 to 0.21 on the center of the Grf and from 0.2 to 0.28 on the edges of the Grf. This increase is partly due to uniaxial pressure employed during the densification step which tend to align the Grf, within the copper matrix, in a plane perpendicular to the pressure direction. Also, additional cause to the degradation of Grf can be due to the difference of coefficient of thermal expansion (CTE) between the Grf and the Cu matrix. Indeed, Cu expands isotropically with a CTE equal to 17.5 ppm.K -1 whereas Graphite expands with a negative CTE (αGr = -1.5 ppm.K -1 ) along their planes and a positive CTE through plane (αGr = 27.1 ppmK -1 ) perpendicular to the planes. This CTE difference may induce stresses and deformation of the Grf during cooling step of the materials (from densification temperature (650 °C to RT).

d) Microstructure of the Cu/Grf composite

Figure 12 shows side view micrographs of Cu/Grf composites with 5, 10, 20, 30 and 40 vol.% of Grf. On these micrographs, light contrast represents the Cu matrix and dark one the Grf reinforcement which exhibits lamellar morphology. During the densification step, the platelet geometry (large length to thickness ratio) of the Carbon reinforcement induces a preferred orientation of Grf in a plane perpendicular to the pressing direction. However, some local disorientations, decreasing with increase of Grf volume fraction from 8,1 ° for 5 vol.% Grf and 6,7 ° for 40 vol.% Grf (measured by image analysis on around 1000 Grf with the ImageJ software), with respect to this plane are shown by red circles on Figure 12. These misoriented regions could be explained by the intrinsic morphology of the dendritic Cu and the contact between Grf. Alignment of flakes above 30 vol.% becomes better and help to increase the TC in plane perpendicular to the pressing direction [START_REF] Firkowska | The Origin of High Thermal Conductivity and Ultralow Thermal Expansion in Copper-Graphite Composites[END_REF]. With the increase of Grf content, distances between flakes are shorter and the composite microstructure is dominated by Grf geometry.

No differences of microstructure were observed between RA and 3D mixing.

e) Effect of the reductive treatment on the composite density Figure 13 shows the evolution of the relative density of Cu/Grf composite materials with Grf volume content. From that figure, it is clear that relative densities decrease when the volume fraction of Grf increases and mostly follow classical rule of mixture. In addition to the porosity, which is inherent to the powder metallurgy route, the decrease in the relative density may be correlated with Cu-Grf interfacial voids due to non-chemical affinity between Cu and Grf and to the Oxygen functional groups present on Grf surface for Set A and B. For Set C, relative densities are higher on the entire concentration range comparing to Set A and B. As demonstrated with TGA (Figure 8) and Raman spectroscopy (Table 4) analysis, the reductive treatment in Set C may induce the shrinkage of Grf. This could be due to the desorption of impurities and interstitial oxygenated groups, and could explain the increase of relative density.

However, relative densities were found to be all above 98 % for all set of materials.

f) Effect of the powder process methodology on thermal conductivity Figure 14 summarized in-plane TC and Oxygen content for a Cu/40 vol.% Grf composite obtained after different powder processing methodology. Each value is an average of measurements on three different samples and errors are no bigger than marker spot. Our experimental values are compared with Cu/Grf materials fabricated using ball milling process in [START_REF] Firkowska | The Origin of High Thermal Conductivity and Ultralow Thermal Expansion in Copper-Graphite Composites[END_REF]. Several points can be pointed out from this figure:

1. Our experimental TC values are all higher than TC value obtained with ball milling process [START_REF] Firkowska | The Origin of High Thermal Conductivity and Ultralow Thermal Expansion in Copper-Graphite Composites[END_REF] 2. TC is higher for material mixed with RA mixer (Set B) than material mixed with 3D mixer (Set A)this behavior can be attributed to the low amount of defects induce by the RA Thereby, the TC out-of-plane decreased with the volume fraction of Grf increases, while inplane TC increased over the entire range of concentration even at low vol.% of Grf (below 20 vol.%). This was not the case in [START_REF] Firkowska | The Origin of High Thermal Conductivity and Ultralow Thermal Expansion in Copper-Graphite Composites[END_REF] where for low content of Grf (with similar particulate size, around 300 µm) the in-plane TC of Cu/Grf composite was below Cu TC. The corresponding TC of this composite is 630 W.m -1 .K -1 which is close to 60 % higher than pure Cu.

Conclusion

Cu/Grf composite materials with different volume fraction of Grf (from 0 to 40 vol.%) have been fabricated via vacuum uniaxial hot-pressing process. Effect of powder processing methodology on C-lattice parameters, Grf quality (defects, oxygenated sp 3 -C), final composite Oxygen content and TC has been investigated. Results could be summarized as-follow:

1) Reductive thermal treatment at 400 °C on Grf improves crystallinity structure as verified by XRD and Raman spectroscopy measurements. Elimination of points defects scattering centers such as impurities, interstitials and surface oxygenated sp 3 -C is taking place during this treatment.

2) A comparison between three mixing processes shows that the RAM device is an efficient mixer for the fabrication of Cu/Grf composite. It allows faster mixing and elimination of defects on Grf edges without powder oxidation. Hence, the decrease of oxygen content in the final composite leads to better relative density.

3) Excellent TC up to 630 W.m -1 .K -1 for Cu/40 vol.% Grf was obtained with this process combining new PPM and densification via vacuum hot-pressing. Even though, the interface between Cu and Grf was not studied in this work, Cu/Grf composite demonstrates better TC than composites fabricated using carbide forming elements [14] [15] or metal coating on graphite [16] [17].

Figures

a) Characterization on Graphite flakes powderFigure 4

 4 Figure4shows the FTIR spectrum of as-received Grf. Native graphite should present only two main peaks around 1670 cm -1 and 1535 cm -1 attributed to C=C aromatic. For our Grf, broad and intense band of O-H stretching vibration appears at 3423 cm -1 . Around 1448, 1385, 1115, 1016 cm -1 bands corresponding to O-H stretching and bending of intercalated water in the

c)

  Effect of mixing processes on Graphite flakes quality Powder mixing (Grf and Cu powder) was performed using a RA mixing and compared with more traditional 3D mixing. Summarized results are presented in Erreur ! Source du renvoi introuvable. (average of at least four acquisitions). Contrary to ball-milling [1], 3D and RA mixing helps to decrease the number of defects on the edges of the Grf. The increase of the D peak on the center of Grf, with 3D mixing, indicates the additional introduction of defects to the C-lattice as well as a decrease of the crystalline size. It has been demonstrated that it is possible to know the motion of mixing by calculating the Froude number Fr from the equation developed by Mellmann in

mixer 3 .

 3 TC is improved after the annealing treatment under reductive atmosphere of the Grf before densification (Set C compare to set B)this behavior can be explained by the improvement of theoretical TC value of the Grf 4. The content of oxygen decreases when the TC of the composite material increases and Cu/Grf material after Grf reductive treatment has an oxygen value almost 6 times lower than Cu/Grf (no treatment) 5. Maximum TC obtained from Set C (630 W.m -1 .K -1 ) is 1.6 time higher than Cu matrix obtained by powder metallurgy process (380 W.m -1 .K -1 ) Erreur ! Source du renvoi introuvable. shows the evolution of the in-plane and out-of-plane TC for the Set C composite material. The anisotropy of TC values of the Cu/Grf composite materials is fully induced by the Grf TC anisotropy and the orientation of the Grf in-plane.

Figure 1 .

 1 Figure 1. SEM micrographs of dendritic Cu powder.

Figure 2 .

 2 Figure 2. SEM micrographs of Grf powder.

Figure 3 .

 3 Figure 3. Definition of in-plane and out-of-plane directions in bulk Cu/Grf composite material. a) CAO model of unpolished samples (Ø 6 × 3.5 mm) obtained after machining dense Cu/Grf composite material (Ø 20 x 8 mm). b) CAO model of a Cu/Grf sample cut in the middle and polished revealing out of plane and in-plane directions of Grf.

Figure 4 .

 4 Figure 4. FTIR spectrum of as-received Grf. Dark blue lines correspond to peaks positions in wavelength (cm -1 ) related to -OH stretching and bending vibrations. Purple lines: positions of stretching and bending of aromatic rings. Light blue lines: positions of epoxy, carbonyl and carboxyl groups. Green lines: positions of CH symmetric and asymmetrical vibrations of methyl and methylene.

Figure 5 .

 5 Figure 5. XPS Survey spectrum of as-received Grf.

Figure 6 .

 6 Figure 6. High resolution core level of a) C1s and b) O1s XPS of as-received Grf.

Figure 7 .

 7 Figure 7. Raman spectra from Grf center a)1 & edge a)2. The micrograph b) shows analysis locations on Grf.

Figure 8 .

 8 Figure 8. TGA analysis of Grf under Ar/H2(5%) reductive atmosphere. Evolution of the weight loss of Grf during elevation of temperature (from 50 °C to 1000 °C).

Figure 9 .

 9 Figure 9. Representative X-Ray Diffraction patterns of Grf obtained after reductive treatment under Ar/H2(5%) at different bearing temperatures. (Peaks between different reductive treatments were shifted in both x and z for easier interpretation).

Figure 10 .

 10 Figure 10. Thermal conductivity vs d(002) of Carbon fibers and ThermalGraph ® panel. Adapted from [10].

Figure 11 .

 11 Figure 11. Point defects (surrounded in red) corresponding to folded edges and broken parts of Grf.

Figure 12 .

 12 Figure 12. SEM micrographs of in-plane Cu/Grf composites with a) 5, b) 10, c) 20, d) 30 and e) 40 vol.% of Grf. Red circles represent location of misorientations of Grf in Cu matrix. Scale bars correspond to 500 µm.

Figure 13 .

 13 Figure 13. Plot of the density of the Cu/vol.% Grf composites for the different sets used, in comparison with the theoretical (expected) density obtained by the rule of mixture.

Figure 14 .

 14 Figure 14. Effect of powder processing methodology on Cu/40 vol.% Grf composite oxygen content and TC perpendicular to the compressive axis at 70°C [1].

Figure 15 .

 15 Figure 15. Thermal conductivity of Cu/Grf composite materials with volume fraction of Grf dependence. In-plane TC is represented by the red line and out-of-plane TC by the blue line.

Table 1 .

 1 XPS qualitative analysis of Grf.

	Tables
	At. % C

Table 2 .

 2 XRD parameters of Grf before and after reductive treatment at different bearing temperatures.

	Temperature of reduction (°C)	(hkl)	d (Å)	FWHM (°)	Grf Kc obtained from [10] (W.m -1 .K -1 )
	Single crystal	(004) (002)	1.678 3.356	-	1600
	As-received Grf	(004) (002)	1.682 3.377	-0.178	500
	200	(004) (002)	1.682 3.377	-0.173	500
	400	(004) (002)	1.679 3.360	-0.132	1130
	800	(004) (002)	1.682 3.373	-0.157	640

Acknowledgments

The authors would like to thank Eric Lebraud, Dominique Denux, and Christine Labrugère of the Institut de Chimie de la Matière Condensée de Bordeaux for performing the XRD, TGA and XPS analyses, respectively.

Declarations of interest: none

Funding: This research did not receive any specific grant from funding from funding agencies in the public, commercial, or not-for-profit sectors.

Table 3. Integrated intensity ratio AD/AG and average G and 2D peak positions of Grf after reduction under Ar/H2(5%).

Table 4. Calculated integrated intensity ratio AD/AG and average G and 2D peak positions of Grf after different type of mixing.

As-received Grf 0,13

Temperature of reduction (°C)

Center Edges