Ph.D Loic Constantin

Yongfeng Lu

Jean-Francois Silvain

3D LASER PRINTING OF DIAMOND-BASED COMPOSITES FOR THERMAL MANAGEMENT APPLICATIONS

The constant increase of the working frequency of semiconductor-based devices with their miniaturization led to severe overheating, which affect their lifetime and reliability.

Hence, thermal management has become a significant concern for the microelectronic area and needs to be addressed. Diamond (D) is known to be an excellent material for thermal dissipation as it possesses one of the highest thermal conductivity (TC) of any natural material and has a high electrical resistivity. D can cool electronic chips in two ways. When used in the form of a film, D acts as a heat spreader. When utilized in powder-form, Ds can be introduced into metals to enhance their TC and bring dimensional stability at elevated temperatures. The resulting metal/D composite materials are thus, excellent component to form heat sinks. Naturally, the thermal performances of heat sinks are closely related to their surface area. Although the attractiveness of D-based materials in term of thermal performance, they often exhibit simple geometry mostly due to the complexity of machining D-based materials into intricated designs. In this study, the laser 3D printing of copper/D composite materials is proposed to fabricate highly complex Cu/D structures which could remodel their applications. Before additively manufactured Cu/D composite materials, several challenges need to be addressed. First, the additive manufacturing of pure Cu is optimized and characterized. Then, due to a lack of a chemical affinity between Cu and D, the Cu-D interfacial zone is introduced in the composite material. Later, a molten salt coating process is studied to produced graded and multilayer coating of oxide/carbide and carbide/carbide, respectively, on carbon materials. Next, the additive manufacturing of highly sophisticated Cu/D composite structures is presented. Finally, the deposition of D films is performed by laser-assisted combustion flame. The effects of introducing ultraviolet lasers into the combustion flame are characterized in terms of chemical reaction and D film quality and growth rate.

IMPRESSION 3D DE MULTI-MATERIAUX

I also would like to thank Dr. Lisha Fan for the numerous advices, help, and time she provided during all the years. Also, I would like to thank Zhipeng Wu, and Nada Kraiem for their contribution to this work. I want to thank Prof. Jean-Luc Battaglia for his contribution and help on the thermal conductivity measurement and Prof. Bai Cui for providing free access to his labs. I want to thank Prof. Jerome Roger for his help and impact to the work.

I want to thank all my dear colleagues from the University of Nebraska Lincoln, Aofei Mao, Xi Wang, Lei Lu, Peixun Fan, Timothy Carlson, Nan Li, and from the ICMCB, Benjamin Thomas, Mael Pontoreau Clio Azina, Adrien Morvan and Nan Chamroune. Also, I would like to thank Pamela Weise, Laurie Prettyman, and Teresa Ryans, which were

always available to answer questions and help.

Finally, I would like to dedicate a special thanks to my family and friends for being present and supportive during this program.

RESUME DES TRAVAUX DE THESE

Les industries et les systèmes de transport reposant sur l'énergie électrique ont connu une croissance rapide depuis les années 1950. Les dispositifs semi-conducteurs représentent le coeur des systèmes électroniques de puissance. Ils permettent de contrôler l'énergie électrique en la distribuant, la générant, la stockant et la convertissant. Ces systèmes se trouvent dans les transports, tels que l'automobile, les chemins de fer et l'aéronautique, et dans les systèmes de production, de conversion et de distribution d'électricité.

Au cours du dernier demi-siècle, des progrès significatifs ont été réalisés pour réduire leurs dimensions du micromètre à l'échelle nanométrique. Une telle miniaturisation permet d'intégrer un grand nombre de puces extrêmement proches les unes des autres, réduisant ainsi la longueur du trajet du signal et contribuant à des fréquences de fonctionnement plus élevées.

Cependant, l'augmentation constante de la puissance électrique et la miniaturisation des appareils ont conduit à des problèmes de gestion thermique dus à des densités de flux de chaleur critiques. Le flux thermique était de 0,5 W/cm [START_REF] Ferain | Multigate transistors as the future of classical metal-oxide-semiconductor field-effect transistors[END_REF] Ensuite, des méthodes de fabrication additive assistées par laser sont présentées, telles que la Selective Laser Sintering (SLS), Selective Laser Melting (SLM) et Laser-Directed

Energy Deposition (LDED). Les principaux défauts induits dans les matériaux au cours de chacun de ces processus sont présentés ainsi que les phénomènes physiques qui leur sont liés. Enfin, les progrès de l'impression laser 3D de structures complexes pour des applications de gestion thermique sont revus.

Au chapitre 3, la fabrication additive de Cu pur par les procédés SLM est décrite. Dans cette étude, des pièces fonctionnelles en Cu sont imprimées. Une enquête sur les paramètres d'impression est effectuée pour optimiser la qualité d'impression et comprendre les mécanismes de formation. Ensuite, des dissipateurs thermiques en Cu avec diverses conceptions structurelles sont fabriqués en utilisant les paramètres d'impression optimisés.

Leurs propriétés physiques et leurs performances thermiques sont étudiées. Les dissipateurs de chaleur présentent une conductivité thermique de 368 W/m.K. Les performances thermiques des dissipateurs thermiques imprimés sont ensuite comparées à celles du commerce en étudiant leur capacité de refroidir une puce électronique. Il est montré que les dissipateurs thermiques imprimés avec des structure complexes surpassent de 45% la dissipation thermique du dissipateur du commerce (Figure 1). Dans le chapitre 6, l'influence de la photolyse induite par laser UV sur la croissance de D est étudiée. L'Optical Emission Spectroscopy (OES) et la Laser-induced Fluorescence (LIF) de la flamme indiquent que la photolyse induite par laser UV produisait de grandes quantités de radicaux réactifs. Ces réactifs contribuent directement à la promotion de la croissance D. L'étude du processus de nucléation suggère que l'irradiation au laser excimère KrF modifie la croissance du diamant de manière à favoriser la formation de xi diamant et à supprimer l'accumulation de carbone amorphe. Cela conduit donc à un taux de dépôt de diamant amélioré et à une qualité de diamant améliorée (Figure 5). De plus, la couche de transition est réduite de 700 nm sans laser à 5 nm avec irradiation laser UV.

Finalement, dans le chapitre 7, les travaux de thèse sont résumé suivi de future axes de recherche. 1 Motivations and Outline

TABLE OF FIGURES AND TABLES

Motivations

Industries and transportation systems relying on electrical energy and power electronics have undergone rapid growth since the 1950s. [START_REF]World Energy Outlook -Topics[END_REF] Acting as electrical energy converters, semiconductor devices represent the heart of power electronics systems. They enable the change of an electrical current from direct to alternating and modification of the potential and the current intensity. Hence, semiconductor devices allow electrical energy to be controlled by distributing, generating, storing, and converting it. These systems are found in transportation, such as automotive, railway, and aeronautics, and in electrical production, conversion, and distribution systems.

In the past half-century, significant progress has been made to reduce devices from microscale to nanoscale; and the aim of future technology is to reach sizes below ten nanometers. [START_REF] Ferain | Multigate transistors as the future of classical metal-oxide-semiconductor field-effect transistors[END_REF] Such miniaturization will permit a large number of chips to be integrated extremely close to each other, thereby reducing the signal path length and contributing to higher operating frequencies. [START_REF] Tong | Advanced Materials for Thermal Management of Electronic Packaging[END_REF][START_REF] Sabryz | Thermal analysis and active cooling management for 3D MPSoCs[END_REF] Hence, increasing the density of transistors and other integrated circuits (ICs) will enhance their performance. In 1958, Gordon E. Moore, the co-founder of Intel, noticed that the number of transistors per IC doubled every two years.

He predicted that the trend would continue for decades. This is known as Moore's law and has been the driving force for miniaturization and the performance improvement of IC devices. 5 In 2003, Moore mentioned that the exponential increase in the component number would eventually reach saturation. 6 On the other hand, the constant increase in the electrical power and miniaturization of devices has led to thermal management issues due to critical heat flux densities. Heat flux was 0.5 W/cm 2 in the 1980s, 40 W/cm 2 in the 2000s, and now often reaches 150 W/cm 2 . [START_REF] Anandan | Thermal management of electronics: A review of literature[END_REF][START_REF] Ellsworth | Chip power density and module cooling technology projections for the current decade[END_REF] Moreover, the temperature of a silicon (Si) chip for example, must be maintained below ~150 C to avoid instability and failure. [START_REF] Huque | Silicon-on-insulator based high-temperature electronics for automotive applications[END_REF] It thus appears that the capacity to cool a device to maintain a relatively low temperature is critical to determining its lifetime, reliability, and efficiency. Indeed, more than half of electronic failures are related to overheating, while 6%, 19%, and 20% are from dust, moisture, and vibration, respectively. [START_REF] Anandan | Thermal management of electronics: A review of literature[END_REF] Consequently, thermal management has become the important limiting factor in further development of power electronics.

In power electronic modules, various substrate materials are attached to the chip to conduct and dissipate the heat. As shown in Figure 1.1-1, the chip is mounted on a ceramic substrate to insulate the module and limit electrical losses. The module is then connected to a copper (Cu) base plate and heat sink allowing the heat produced to dissipate. [START_REF] Guillemet | Diamond-based heat spreaders for power electronic packaging applications[END_REF] The heat dissipation capability of the package is restricted to its thermal resistance, in other words, the sum of the thermal resistance of each component. [START_REF] Azina | Diamond-based multimaterials for thermal management applications[END_REF] Thus, the thermophysical properties of each part need to be carefully analyzed to preserve their reliability. A significant concern is the formation of thermomechanical stress between two adjacent layers due to thermal expansion mismatches (coefficients of thermal expansion (CTEs)). If a significant CTE difference between two materials is introduced, tensile or compressive stresses are generated upon thermal cycling, affecting the mechanical integrity of the package and, potentially, eventual failure. 12 Cooling can be divided into two main categories: passive and active. A passive cooling system consists of a natural convection and conduction process. On the other hand, active cooling involves forced convection. [START_REF] Anandan | Thermal management of electronics: A review of literature[END_REF][START_REF] Lee | Optimum design and selection of heat sinks[END_REF] Cooling devices are usually made of highly thermally conductive materials with tailored CTEs. Copper metal matrix composites (MMCs) are among the most advanced materials for thermal management applications.

Diamond (D) is often used as a reinforcement in the MMCs due to its high thermal coefficient (TC) (up to 2000 W/m.K) and low CTE (0.8 x10 -6 /K). [START_REF] Silvain | A review of processing of Cu/C base plate composites for interfacial control and improved properties[END_REF] Copper/diamond composite materials have demonstrated exceptionally high TC with low CTE and show promise for applications in thermal management. [START_REF] Guillemet | Diamond-based heat spreaders for power electronic packaging applications[END_REF][START_REF] Azina | Diamond-based multimaterials for thermal management applications[END_REF] However, the presence of D particles makes it difficult to machine these materials into complex geometries withdraw their applications as heat sinks. Additionally, D is an excellent candidate for use as an insulating layer in electronic packaging due to its unique thermal properties and high electrical resistivity. [START_REF] Asmussen | Diamond Films Handbook[END_REF] Often D films are grown by chemical vapor deposition (CVD), but either the growth rate or crystallinity of the film is altered. [START_REF] Fan | Diamond film deposition using laser-assisted combustion flames[END_REF] To answer these challenges, we proposed using three-dimensional (3D) laser printing to fabricate intricate Cu/D composite designs. In addition, laser-assisted CVD of D film is intended to grow high-quality D films in a relatively short time frame. This dissertation focuses on the fabrication of D-based heat sinks and heat spreader materials. The following topics are covered: 1) additive manufacturing of Cu and Cu/D composites and 2) ultraviolet (UV) laser-assisted D film deposition.

Overview of the dual Ph.D. program

This thesis results from an international collaboration in the framework of a dual Ph.D.

Dissertation outline

The most significant contributions of this dissertation carried out at the LANE lab at UNL and ICMCB at UBX are:

1. Development of a selective laser sintering (SLS) technology to print dense, crack-free Cu structures.

2. Fabrication of complex Cu heat sink structures via selective laser melting (SLM) with high heat-dissipation performance.

3. Development of an innovative coating process for carbon-based materials.

4. Investigation of the kinetics of a molten-salt process to form carbide coatings.

5. Synthesis of a new titanium (Ti)-based coating for Cu/D composite via a molten-salt process. The dissertation includes seven chapters. In Chapter 1, the motivations and the dissertation outline are presented. In Chapter 2, the state of the art related to this work and background information is reviewed. In Chapter 3, the additive manufacturing of pure Cu by the SLS and SLM processes is described, including the effect of the printing parameters.

In Chapter 4, the development of a Ti-based coating on carbon fibers and D particles is shown as well as the synthesis of multilayer carbide coating on carbon fibers. In Chapter 5, LDED and SLM of Cu/D composite materials is presented along with the fabrication of intricate designs. In Chapter 6, the deposition of diamond films by laser-assisted CVD is presented with an emphasis on the effect of the UV laser wavelength on D growth. Finally, Chapter 7 is a summary of the work followed by future research directions.

State of the art

Introduction

The constant increase in power and packing densities in power electronic devices has led to heat dissipation issues. Significant demand for developing efficient heat-dissipating materials with high TC and a low CTE has appeared over the past decades. [START_REF] Luedtke | Thermal Management Materials for High-Performance Applications[END_REF]

Synthesis of diamond and properties

Diamond has fascinated humankind for thousands of years because of its stunning beauty and was first used as a gemstone. Later, a growing interest among the scientific community appeared due to its outstanding properties. Diamond is a unique material in the sense that it combines extreme hardness and wear resistance, high thermal conductivity and electrical resistivity, low coefficient of thermal expansion, chemical inertness, and high optical transparency. Its excellent properties make it attractive for numerous applications, such as drilling and cutting, optics, heat sinks, and electronic packaging. 20

Diamond synthesis

Catalytic high-pressure and high-temperature diamond

The catalytic high-pressure and high-temperature (HPHT) approach is the most advanced method of producing a large quantity of synthetic D powders. This method consists of applying high pressure (5 to 7 GPa) and temperature (1300 to 2000 ºC) to highpurity graphite powders. The graphite-to-diamond conversion is assured by adding transition metals (e.g., iron (Fe), nickel (Ni), cobalt (Co), manganese (Mn)), which serve as solvents and catalyzers. At high temperature, the molten metal dissolves the graphite;

then, under high pressure, homogenous nucleation of D particles occurs. [START_REF] Sung | Graphite → diamond transition under high pressure: A kinetics approach[END_REF] The growth of nuclei into D powders is then governed by the diffusion of C atoms across the molten metal.

Often, cubic, cubo-octahedral, or octahedral D are obtained. Their sizes can go from micrometric up to centimeter depending on the time and conditions of the HPHT process. 20

Chemical vapor deposition

Chemical vapor deposition differs from the HPHT method mainly because D is synthesized at low pressure (atmospheric) and low temperature (> 1000 ºC). In the CVD method, the reaction is driven by a chemical reaction in a gas phase. The nucleation and growth of D occurs on a substrate. [START_REF] Asmussen | Diamond Films Handbook[END_REF] The gas mixture is usually composed of hydrogen (H), oxygen (O), and hydrocarbons (e.g., ethylene (C2H4), acetylene (C2H2)). [START_REF] Matsumoto | Growth of diamond particles from methane-hydrogen gas[END_REF] The decomposition of hydrocarbons into reactive species (e.g., -CH2-) allows the epitaxial growth of sp 3 and sp 2 carbon species. Later, H and O etch the nondiamond-bonded species and eliminate the incorporation of impurities. [START_REF] Angus | Growth of Diamond Seed Crystals by Vapor Deposition[END_REF] Many CVD systems have been developed over the past decades, such as hot filament, plasma jet assisted, microwave, and combustion flame. [START_REF] Fan | Ultraviolet laser photolysis of hydrocarbons for nondiamond carbon suppression in chemical vapor deposition of diamond films[END_REF][START_REF] Murakawa | Diamond coating of a long substrate by use of a combustion flame[END_REF] The CVD deposition of D is a competition between growth rate and D quality. Among the CVD systems, combustion synthesis was studied extensively due to its fast growth rate (up to 60 μm/h), as illustrated in Figure 2.2-1(a). [START_REF] Murakawa | Diamond coating of a long substrate by use of a combustion flame[END_REF] In this process, molecules undergo different excitations, such as translations, rotations vibrations, and electronic excitations. Laser-assisted combustion synthesis was introduced to enhance D quality while maintaining or improving the growth rate. Molecular vibrations are in the infrared (IR) range; and hence, by shining an IR laser through the torch, vibrational energy can be selectively provided to reactive species. The laser-assisted combustion of D was demonstrated using a carbon dioxide (CO2) laser. It was shown that the vibrational excitation of C2H4 can be achieved in a combustion flame, as shown in Figure 2.2-1(b). [START_REF] Xie | Resonant excitation of ethylene molecules in the combustion flame CVD of diamond using a wavelength tunable CO2 laser[END_REF][START_REF] Xie | Control of crystallographic orientation in diamond synthesis through laser resonant vibrational excitation of precursor molecules[END_REF][START_REF] Fan | Influence of Laser Vibrational Excitations of Ethylene Molecules in Laser-Assisted Combustion Diamond Synthesis[END_REF][START_REF] Fan | Mass spectrometric investigation of the roles of several chemical intermediates in diamond synthesis[END_REF] The laser wavelength needs to precisely match the vibrational modes of molecules i.e., 10.532 μm for C2H4. In addition, the D growth rate and quality were both improved. Notably, a 10% enhancement of the D quality was reported with a growth rate three times faster. [START_REF] Guillemet | Diamond-based heat spreaders for power electronic packaging applications[END_REF][START_REF] Azina | Diamond-based multimaterials for thermal management applications[END_REF]

Mechanical and thermal properties

Mechanical properties

It is well known that diamond is the hardest natural material with a ten on the Mohs hardness scale (1 to 10). This property is essential due to its high atomic packing density (1.77 x 10 23 atoms/cm 3 ). This stunning hardness makes it an excellent material as a protective coating against wear. Also, D has a high Young's modulus with a value of about 1050 GPa, approximatively fives time superior to that of steel. Its high modulus also signifies that it has low compressibility and a Poisson's ratio of 0.07. 20

Thermal properties

In addition to its mechanical properties, D has the highest thermal conductivity of any natural material, making it the predominant material in thermal management applications.

The thermal conductivity of high-purity D can go as high as 2200 W/m.K at room temperature, which is more than five times that of pure Cu (400 W/m.K). This remarkably high thermal conductivity results from strong covalent bonding and low phonon scattering.

However, this value was measured on a pure D crystal. Synthetic D films exhibit thermal conductivities between 300 and 700 W/m.K (Figure 2.2-2). [START_REF] Takabatake | Thermal characterization of CVD diamond film by photoacoustic method[END_REF][START_REF] Ho | Measurement of thermal conductivity in diamond films using a simple scanning thermocouple technique[END_REF][START_REF] Wolter | Thermal conductivity of epitaxially textured diamond films[END_REF] The lower thermal conductivity of D films compared to the pure single crystal is essentially due to the presence of impurities, grain boundaries (polycrystalline films), and crystallographic defects (dislocations, twin boundaries). Phonons carry heat in the D crystal lattice and are scattered in multiple directions due to the presence of defects (i.e., impurities, grain boundaries, etc.), resulting in lower thermal conductivity. [START_REF] Anthony | The thermal conductivity of isotopically enriched polycrystalline diamond films[END_REF] Note that most of the D films obtained via CVD show a black color, indicating the presence of nondiamond C.

Nonetheless, CVD diamonds, with their preferred crystallographic orientation, allow the phonon-free path inside the film structures to be enhanced, making them valuable for thermal management applications in electronics. Furthermore, D has the lowest heat capacity in nature with a value of 6.19 J/mol.K in a temperature range of ~0 to 800 K, about four times smaller than Cu (24.2 J/mol.K). The low heat capacity of CVD diamond makes it a perfect material for a heat spreader. In other words, a heat sink needs to sustain a large amount of heat while maintaining its temperature (i.e., high heat capacity). [START_REF] Asmussen | Diamond Films Handbook[END_REF] However, D has low heat capacity and high thermal conductivity, thus it will spread the heat efficiently when combined with a heat sink.

Finally, D presents a low CTE, from 0.8 x 10 -6 to 4 x 10 -6 /K in a temperature range from 300 to 1200 K. The low CTE makes D a desirable material not only for heat spreaders on Si chips (CTE = 2.6 x 10 -6 /K) but also as reinforcement in MMCs. In general, metals have high CTEs (< 10 x 10 -6 /K), causing the metal heat sink to peel off from electronic chips after multiple cycles (on and off cycle, heating, and cooling). Thus, using D as a reinforcement could lower the CTEs of MMCs and increase their thermal conductivity. 20

Copper/diamond composites

Copper provides an efficient way to remove heat in microelectronics because of its high TC. However, the CTE mismatches between Cu (17 x 10 -6 /K ) and ceramics substrates (4-9 x 10 -6 /K) leads to delamination, which causes electronic components to overheat. The fabrication of components with high TC, low CTE, and intricate shapes is thus crucial for the development of highly efficient heat-exchanging components for microelectronic industries. [START_REF] Vasiliev | Heat pipes in modern heat exchangers[END_REF][START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF] Creating a composite material consisting of two highly thermally conductive elements, such as Cu and D, may fulfill these unmet needs. [START_REF] Qu | Review of metal matrix composites with high thermal conductivity for thermal management applications[END_REF] However, the lack of chemical affinity between Cu and D, together with the difficulty of mechanically processing D-based materials into complex shapes, presents significant challenges. [START_REF] López | The solubility of C in solid Cu[END_REF][START_REF] Silvain | Copper-Carbon and Aluminum-Carbon Composites Fabricated by Powder Metallurgy Processes[END_REF]

Cu/D, a nonreactive system

A reactive system is defined as the assembly of two chemically active phases. In other words, the two components can react with each other. However, Cu and C are a nonreactive system. This means that there is no chemical affinity and almost no solubility leading to a high wetting angle (~145º). [START_REF] Kong | Enhanced wetting and properties of carbon/carbon-Cu composites with Cr 3 C 2 coatings by Cr-solution immersion method[END_REF] Cu. [START_REF] Mortimer | The wetting of carbon by copper and copper alloys[END_REF] Notably, the alloying element reacted with the C substrate and formed a carbide layer, permitting the formation of chemical bonds. [START_REF] Mortimer | The wetting of carbon by copper and copper alloys[END_REF] A. Veillère et al. demonstrated that the addition of chromium (Cr) into a Cu and C powder mixture led to the creation of chromium carbide (Cr3C2) at the interphase after an annealing step. [START_REF] Veillère | Relationship between interphase chemistry and mechanical properties at the scale of micron in Cu-Cr/CF composite[END_REF] Others reported on liquid sintering, which consists of adding an alloying element having a lower melting point than Cu, and thus turning it into a liquid during the sintering process. [START_REF] Chung | Effect of Titanium Addition on the Thermal Properties of Diamond/Cu-Ti Composites Fabricated by Pressureless Liquid-Phase Sintering Technique[END_REF] The liquid-phase sintering of Cu/C composites containing a small amount of alloying element, molybdenum (Mo), tungsten (W), or Ti, showed the formation of carbide interphases between the Cu and C reinforcement. [START_REF] Chung | Effect of Titanium Addition on the Thermal Properties of Diamond/Cu-Ti Composites Fabricated by Pressureless Liquid-Phase Sintering Technique[END_REF] However, the properties were severely degraded due to inhomogeneous dispersion of the alloying elements at the interfaces. In addition, traces of alloying elements remained in the Cu matrix, resulting in a lower TC compared to that of pure Cu.

Ex situ formation of interphases for Cu/D composites

Another approach is to coat the D or C reinforcements before the sintering step. Multiple processes were studied, such as sputtering, microdeposition, or molten-salt synthesis. Liu [START_REF] Kang | Preparation of copper-diamond composites with chromium carbide coatings on diamond particles for heat sink applications[END_REF] Interphase is thus vital to achieving highly thermally conductive Cu/D composites.

Fabrication of Cu/D composites

A large variety of processes have been used to fabricate MMCs, which can be divided into solid and liquid manufacturing. The conventional fabrication of MMCs uses powder metallurgy, which is a solid process. In this method, the densification is realized by sintering. Multiple metallurgy methods have been developed, depending on the heating sources, such as hot pressing ((HP), induction coil), SPS (electric current in the material), or microwave sintering (microwave). In all of these methods, the powder mixtures are densified at a temperature lower than the melting point and under high pressure (uniaxial), as illustrated in Metal matrix composites can also be manufactured via liquid processes, such as stir casting or infiltration. In stir casting, the metal powders containing the reinforcement are totally melted and poured into a mold. However, due to the density difference between the two components, heterogeneous dispersion of the reinforcement may occur. In the infiltration process, a preform made of reinforcement is fabricated first using binders. Then a liquid metal is infiltrated into the preform via capillarity force and/or pressure.

Nevertheless, contamination from the binders may remain in the materials. Once again, the shape of the MMCs is limited by the shape of the mold. Hence, the manufacturing of complex shapes remains challenging; and a new process needs to be developed. [START_REF] Silvain | A review of processing of Cu/C base plate composites for interfacial control and improved properties[END_REF][START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF] 2.4 Laser-assisted additive manufacturing: an overview

Additive manufacturing (AM), or 3D printing, was conceptualized back in the 1970s by David E.H. Jones in a paper published in the New Scientist journal and quickly developed for polymer materials. 49 In the 1980s, the concept of using a laser as a heat source to fuse powders was introduced on polymer powders. But it was only in the late 1990s that the research on AM of metals via laser sintering and melting started. Later in the 2000s, 3D printing of metals matured; and commercial printers became available. [START_REF] Sames | The metallurgy and processing science of metal additive manufacturing[END_REF] Since the start of 3D printing back in the 1970s, a large selection of materials have been 3D printed (polymers, metals, ceramics, and composites) using different technologies (selective laser melting or sintering, direct energy deposition, etc.).

Selective laser sintering/melting

Selective laser sintering or melting refers to a powder bed technology. In this technique, the powders are added layer upon layer, as illustrated in Figure 2.4-1. [START_REF] Sames | The metallurgy and processing science of metal additive manufacturing[END_REF] First, a thin powder layer is deposited onto the working area. Then, the laser selectively melts or sinters the powders into the designed shape. Next, the build plate is lowered; and a new layer is deposited. By repeating the process several times, complex 3D objects can be manufactured. Also, the printing chamber is filled with shielding gas to prevent oxidation of the metallic powders during the laser-matter interactions. The design of the 3D part is done using CAD software, which is then sliced into a print file and loaded into the printer. In SLS, the laser beam lasts a few milliseconds on powder particles, which implies that the sintering time is too short to obtain a dense material. [START_REF] Kruth | Binding mechanisms in selective laser sintering and selective laser melting[END_REF][START_REF] Tolochko | Balling processes during selective laser treatment of powders[END_REF] One solution is to decrease the particle size, which leads to a faster sintering rate (i.e., lowering the surface area is the driving force of sintering). Also, small particles provide higher surface areas and increase laser absorption energy, working temperature, and, thus, sintering kinetics. However, when the size is too small, particles form agglomerates; and the high surface-to-volume ratio is lost. [START_REF] Kingery | Densification during Sintering in the Presence of a Liquid Phase. I. Theory[END_REF] Another solution to overcome this issue is liquid phase sintering. In this sintering process, a second component (called a binder) with a lower melting point than the main component is added in the powder bed. Due to the lower melting point, the binder liquifies during the pass of the laser. Then metal powders are embedded in the binder, and a dense part is obtained. [START_REF] Li | Selective laser melting W-10 wt.% Cu composite powders[END_REF] The densification by liquid sintering can be divided into three steps: 1) when the liquid phase appears, there is a rapid bonding due to the capillary forces exerted on the liquid, called the rearrangement; 2) densification by solution reprecipitation, and 3) solid sintering which causes further densification. Steps 2 and 3 occur by atomic diffusion and, thus, require time to occur. In the SLS process, due to a short laser-matter interaction time, only the rearrangement plays a significant role. During laser sintering, the liquid flows and fills the pores between the metal particles and cools down to form a solid. [START_REF] Anestiev | Model of the primary rearrangement processes at liquid phase sintering and selective laser sintering due to biparticle interactions[END_REF] In SLM, however, the metal powders melt to form a transit pool of the molten material which induces several phenomena. [START_REF] Yap | Review of selective laser melting: Materials and applications[END_REF] During the laser-matter interactions, a spheroidal molten pool is created in a few milliseconds. The presence of liquid metal induces the presence of a plasma jet or plume (i.e., metal vapor) above the molten pool. This gas generates a recoil pressure on the molten pool, which can induce spatter ejection. In addition, a temperature gradient forms inside the pool due to the temperature gradient between the center and the edges. A metal flow from the hot to the cold area arises in the pool, called the Marangoni effect. [START_REF] Wang | Mechanisms and characteristics of spatter generation in SLM processing and its effect on the properties[END_REF][START_REF] Ly | Metal vapor micro-jet controls material redistribution in laser powder bed fusion additive manufacturing[END_REF] The pool dynamic is thus a significant parameter when considering the printing of dense and smooth structures. In addition, the metal viscosity and surface tension of the liquid metal with the surrounding atmosphere and metal parts plays a vital role in the pool stability and is discussed in the next section. [START_REF] Matthews | Denudation of metal powder layers in laser powder bed fusion processes[END_REF][START_REF] Qiu | On the role of melt flow into the surface structure and porosity development during selective laser melting[END_REF] Printing defects are the major limitation of 3D laser printing technologies (i.e., SLS and SLM) as they degrade the properties of the printed part. When the laser strikes the powders, it locally heats them and induces the creation of a thermal gradient. The top layers (hot) tend to expand, while the bottom layers (colder) restrict this expansion. [START_REF] Brown | Fundamentals of Laser-Material Interaction and Application to Multiscale Surface Modification[END_REF] This results in compressive stresses on the top layer and may cause plastic deformation. Later, when the material cools down, those stresses are converted into tensile stresses and may induce cracks. [START_REF] Ali | Effect of scanning strategies on residual stress and mechanical properties of Selective Laser Melted Ti6Al4V[END_REF] Varying the laser scan strategy plays a crucial role in the distribution of heat and thermal stresses. [START_REF] Sames | The metallurgy and processing science of metal additive manufacturing[END_REF] Balling is ubiquitous in SLM and is due to insufficient wettability with the previous layers and to the molten metal surface tension. The formation of ball-like metallic powders can aggregate in subsequent layers and cause defects on the powder bed. [START_REF] Yap | Review of selective laser melting: Materials and applications[END_REF] Finally, the porosity is one of the most critical issues, both in SLM and SLS. Pores often lead to a degradation in physical properties. They are classed into two categories: 1) powder-induced-during atomization of the powders, gas can be trapped into them and released during the process, leading to the formation of spherical pores; and 2) processinduced-if the applied energy is not enough to melt the metal (lack of fusion) or if the energy is too high, spatter ejection phenomena appear (keyhole formation). [START_REF] King | Observation of keyhole-mode laser melting in laser powder-bed fusion additive manufacturing[END_REF] The formation of process-induced porosity is related to the printing parameters. Thus, it is essential to optimize the printing parameters and understand how defects are formed. By using an optimal set of parameters, the density of the material can reach 99% of the theoretical density. [START_REF] Carter | Laser Powder Bed Fabrication of Nickel-Base Superalloys: Influence of Parameters; Characterisation, Quantification and Mitigation of Cracking[END_REF][START_REF] Kruth | Selective laser melting of iron-based powder[END_REF]

Laser-directed energy deposition

Another common laser-assisted AM is LDED. In this AM technology, the laser melts the powders as in SLM; but the part is built by melting a surface (build plate) and simultaneously applying the metal powder, as shown in Figure 2.4-2. The powders are blown through a nozzle, encounter the laser, melt, and deposit. [START_REF] Herzog | Additive manufacturing of metals[END_REF] Cracks, porosity, and balling are the significant defects in LDED, similar to SLM. The advantage of LDED over powder-bed technology is that LDED provides a faster building rate and larger building volumes. However, the surface finish is not as neat as that in powder-bed technology; and post-AM processing is often necessary. This technology is commonly used to coat and repair parts via cladding. 67

Wettability and viscosity

Wetting is a crucial aspect of the AM of metal parts. The wettability of a solid by a liquid is related to the surface tension of the liquid-solid, solid, vapor, and liquid-vapor interfaces. Wettability is defined by a contact angle θ, as can be seen in Figure 2.4-3, where an angle higher than 90° indicates a nonwetting behavior and lower than 90° indicates good wetting between the components. Wettability is expressed as a function of the surface tension of each phase by the following equation. 68

cos θ = γ SV -γ SL γ LV , (1)
where γSV, γSL, and γLV are the surface tension between solid-vapor, solid-liquid, and liquidvapor, respectively. The surface tension depends on the temperature and can be expressed, in the case of Cu powders, as 69

γ = 1330 -0.23(T -1085), (2)
In SLM and LDED, the wettability between a solid and a liquid is influenced by the viscosity (i.e., temperature), the shielding gas, and the presence of impurities. [START_REF] Agarwala | Direct selective laser sintering of metals[END_REF] The viscosity of the molten pool is another significant concern in AM. The viscosity of the liquid phase must be low enough for it to flow and fuse with the surrounding particles but high enough to avoid the balling effect (i.e., surface tension). The viscosity of molten metal can be expressed by 70

μ = μ 0 (1 - 1-φ L φ m ), ( 3
)
where µ0 is the base viscosity and includes a temperature term, φ L is the volume fraction of the liquid phase, and φ m is the critical volume fraction of solids above which the mixture has an infinite viscosity. If the metal is totally liquid, the dynamic viscosity of the liquid is defined by the following equation.

where m is the atomic mass, k is the Boltzmann constant, T is the temperature, and γ is the surface tension of the liquid. As can be seen in the previous equations, the viscosity decreases when the working temperature increases; thus, rheological properties are enhanced. During 3D laser printing, the laser power, scan speed, and layer thickness will influence the temperature of the molten metal and thus affect the rheological properties.

Laser absorption by metallic powders

Laser-matter interactions are an essential parameter in AM for the determination of the processing parameters, such as laser wavelength, power, and scan speed. Two primary mechanisms of light absorption are possible in the case of metal powders. The first is called multiphoton ionization and involves the simultaneous absorption of multiple photons by electrons, which then are ejected from the conduction band (wavelength lower than 1 μm).

The second mechanism is the absorption of laser radiation by free electrons called the photothermal process through inverse bremsstrahlung (e -(slow) + hν → e -(fast)). [START_REF] Chichkov | Femtosecond, picosecond and nanosecond laser ablation of solids[END_REF] For most metals, the thermalization time (time to transfer energy to phonons) ranges from 10 - 12 to 10 -10 s. Thus, if the laser-induced excitation rate is lower than the thermalization rate, the laser energy absorbed is directly transformed into heat. In addition, the laser absorption of metals decreases when the wavelength increases but increases with the temperature. [START_REF] Tolochko | Absorptance of powder materials suitable for laser sintering[END_REF] The mixing law can estimate laser absorption in a powder mixture with an error rate of 10% using the absorbance of each component and its volume fraction, as expressed in the following equation.

A = A 1 . γ 1 + A 2 . γ 2 , (5)
where Ai and γi are, respectively, the absorptance and the volume fraction of component i, respectively.

Additive manufacturing of complex structures for thermal management applications

The ability to manufacture highly complex shapes is revolutionizing many industries, such as the defense, aerospace, and medical industries. [START_REF] Hojjatzadeh | Pore elimination mechanisms during 3D printing of metals[END_REF] Much research has been conducted on the fabrication of thin, sophisticated lattice structures made of Ti and aluminum (Al) alloys and steel, [START_REF] Leary | Selective laser melting (SLM) of AlSi12Mg lattice structures[END_REF][START_REF] Li | Influence of cell shape on mechanical properties of Ti-6Al-4V meshes fabricated by electron beam melting method[END_REF][START_REF] Wang | Study on the designing rules and processability of porous structure based on selective laser melting (SLM)[END_REF][START_REF] Wong | Saturated pool boiling enhancement using porous lattice structures produced by Selective Laser Melting[END_REF] as shown in Figure 2.4-4. Lattice structures are studied extensively for their unique mechanical, thermal, and acoustic properties but also for their lightweight and high energy absorption. [START_REF] Jafari | The utilization of selective laser melting technology on heat transfer devices for thermal energy conversion applications: A review[END_REF] This type of structure has applications in numerous areas and presents significant advantages for heat dissipation applications (i.e., high surface aspect). Few reports study the potential of using SLM to manufacture complex heat dissipaters, as shown in Figure 2.4-5. Heat sinks with different arrays and lattice structures were printed. [START_REF] Wong | Convective heat transfer and pressure losses across novel heat sinks fabricated by Selective Laser Melting[END_REF] The high surface ratio provided by the intricate lattice structure enhances the thermal dissipation performance compared to commercial types of array heat sinks. M.

Wong et al. showed that a diamond lattice structure heat sink enhances thermal dissipation by more than 60% compared to an array heat sink. [START_REF] Wong | Pressure Loss and Heat Transfer Through Heat Sinks Produced by Selective Laser Melting[END_REF] Yet, the printed parts are made of low thermal conductivity metals and alloys, such as steel, Ti, and Al alloys. [START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF][START_REF] Wong | Convective heat transfer and pressure losses across novel heat sinks fabricated by Selective Laser Melting[END_REF][START_REF] Ho | Convective heat transfer performance of airfoil heat sinks fabricated by selective laser melting[END_REF] The additive manufacturing of high TC materials presents several challenges due to their excellent thermal dissipation and often low laser absorption in the near-infrared (NIR) domain. There are difficulties with laser-assisted AM of Cu due to its high thermal conductivity and low laser absorption in the IR domain (> 5 % at λ = 1 μm). [START_REF] Tran | 3D Printing of Highly Pure Copper[END_REF] Over the last few years, attempts were undertaken to 3D laser print Cu parts. Reports show that high laser power is necessary to densify the powder during the SLM process. [START_REF] Colopi | Selective laser melting of pure Cu with a 1 kW single mode fiber laser[END_REF][START_REF] Ikeshoji | Selective Laser Melting of Pure Copper[END_REF][START_REF] Jadhav | Influence of selective laser melting process parameters on texture evolution in pure copper[END_REF][START_REF] Jadhav | Influence of Carbon Nanoparticle Addition (and Impurities) on Selective Laser Melting of Pure Copper[END_REF] reported the manufacturing of dense Cu parts (98%) using a laser power of 800 W (λ = 1080 nm). Yet, they only demonstrated SLM of a simple geometry (i.e., cube). [START_REF] Jadhav | Influence of selective laser melting process parameters on texture evolution in pure copper[END_REF] In addition, a part of the laser power was reflected in the laser optics and caused severe damage. Axel Hess et al., studied the 3D laser printing of Cu using two sets of lasers, a green laser (λ = 515 nm) to heat and thus increase the laser absorption of Cu, followed by an IR laser (λ = 1030 nm) which melted the powders. [START_REF] Hess | Continuous Wave Laser Welding of Copper with Combined Beams at Wavelengths of 1030nm and of 515nm[END_REF] They showed that the use of a green laser improved densification compared to a single IR laser. However, they did not demonstrate the

Conclusions

The objective of manufacturing a heat dissipater for microelectronics is to obtain a material with a low CTE and high TC. Diamond-based materials are an excellent choice due to their unique thermal properties as a polycrystalline film or reinforcement in a metal matrix.

Diamond CVD has a low CTE and heat capacity and high TC. Thus, it is an excellent material for use as a heat spreader when combined with a heat sink. Also, the simplicity, fast growth rate, and cost-effectiveness of using a combustion flame to synthesize D films make it an ideal method for producing heat spreaders. However, the presence of defects during the D growth degrades the thermal properties of the films. Using a laser-assisted combustion flame to grow D appears to be a promising way to improve D quality while maintaining a fast growth rate.

It has been shown that the absence of chemical affinity between Cu and D results in weak and porous interfacial bonds, leading to poor thermal transfer. The importance of having robust interfacial bonding plays a crucial role in the overall thermal properties of the composites. Therefore, it is necessary to produce a chemical bond between the Cu and D reinforcement through interphase. Different carbide interphases and synthesis methods were reviewed, such as MO2C, TiC, and Cr3C2, and made in situ and ex situ. In addition, coating the C reinforcement before the Cu/D manufacturing allows homogenously coated layers to be produced and avoids the inclusion of impurities in the Cu matrix. The manufacturing method of producing MMCs was discussed, such as solid and liquid syntheses. However, it was shown that only simple MMC geometries have been built; and post-processing steps are necessary to fabricate intricately designed structures.

Laser-assisted AM methods were presented, such as SLS, SLM, and LDED. The primary defects induced during each of these processes were presented (i.e., cracks, porosity, and balling) as well as the physics phenomena related to them (i.e., thermal gradient, wettability, and viscosity). Finally, progress on the 3D laser printing of complex structures for thermal management applications was presented. It was shown that intricate structures are allowed to increase the surface ratio and thus promote heat dissipation performance but are currently limited to low thermal conductivity materials.

3 Additive Manufacturing of Pure Copper

Introduction

The high thermal conductivity (TC) (~ 400 W/m.K) of pure Cu makes it an essential material for thermal management applications in numerous industries, such as microelectronics, power plants, and aerospace industry. [START_REF] Silvain | A review of processing of Cu/C base plate composites for interfacial control and improved properties[END_REF][START_REF] Vasiliev | Heat pipes in modern heat exchangers[END_REF][START_REF] Singer | Additively manufactured copper components and composite structures for thermal management applications[END_REF][START_REF] Körner | Additive manufacturing of metallic components by selective electron beam meltinga review[END_REF][START_REF] Lee | Two-phase flow in high-heat-flux micro-channel heat sink for refrigeration cooling applications: Part I--pressure drop characteristics[END_REF] However, the shape (i.e., surface area) of heat sinks plays a crucial role in heat dissipation, and is often a simple geometry due to the current fabrication limitation. [START_REF] Zhang | Experimental study on heat transfer performance of lotus-type porous copper heat sink[END_REF][START_REF] Frazier | Metal Additive Manufacturing: A Review[END_REF] Recently, the development of additive manufacturing (AM), such a selective laser sintering or melting (SLS or SLM), has attracted tremendous attention due to its ability to efficiently manufacture geometrically complex metal parts while minimizing wastes and post-processing steps. [START_REF] Jadhav | Influence of Carbon Nanoparticle Addition (and Impurities) on Selective Laser Melting of Pure Copper[END_REF][START_REF] Frazier | Metal Additive Manufacturing: A Review[END_REF] The thin and sophisticated features enabled by AM and the excellent properties of bulk Cu have the potential to significantly enhance the performance of heat sinks components. [START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF][START_REF] Ramirez | Open-cellular copper structures fabricated by additive manufacturing using electron beam melting[END_REF] SLS of metal-based powders uses a laser as the heat source to heat the powders above the softening temperature. In the conventional sintering (CS) process, bonding occurs when the metal particles form interlinked necks through atomic diffusion at elevated temperatures. [START_REF] Exner | Sintering Processes[END_REF][START_REF] Alexander | The mechanism of sintering of copper[END_REF] However, solid-state diffusion is a slow process, which is difficult to initiate during SLS due to the very fast heating rates and short sintering times (~milliseconds). [START_REF] Johnson | Solid-State Sintering[END_REF] To accelerate the binding process, liquid sintering is widely used in SLS by adding a binder that turns to liquid during the sintering. [START_REF] Agarwala | Direct selective laser sintering of metals[END_REF][START_REF] Tang | Direct laser sintering of a copper-based alloy for creating threedimensional metal parts[END_REF] Liquid sintering occurs as a result of liquid flow and rearrangement. 100 Many research groups have reported the manufacturing of Cu-based metal parts by SLS using various systems (copper-tin (Cu-Sn), copper-iron (Cu-Fe), copper-tin-nickel (Cu-Sn-Ni)) with properties comparable to those obtained by CS. [101][102][103] Although fast sintering for fabrication of metals or alloys can be achieved by adding a binder to the powder mixtures, 104-106 sintering of a single metal remains a challenge due to the necessity of using a large quantity of binders. 107

More recently, attempts have been undertaken to additively manufacture pure Cu via SLM. 82 Silbernagel et al. and SD Jadlhav et al. studied the fabrication of Cu parts by SLM using low laser powers of 200 and 300 W, respectively. [START_REF] Jadhav | Influence of selective laser melting process parameters on texture evolution in pure copper[END_REF]108 Low-density parts (86%) with low TC of 202 W/m.K were reported. The low-density parts resulted from insufficient laser energy deposited on powders due to the fast heat dissipation rate and the low laser absorption by Cu in the near-infrared (IR) range . [START_REF] Hess | Continuous Wave Laser Welding of Copper with Combined Beams at Wavelengths of 1030nm and of 515nm[END_REF]109 A high laser power (800 W) was employed to overcome this issue, and dense Cu parts were printed (97%) with a TC of 336 W/m.K. [START_REF] Colopi | Selective laser melting of pure Cu with a 1 kW single mode fiber laser[END_REF][START_REF] Ikeshoji | Selective Laser Melting of Pure Copper[END_REF] Nevertheless, the high reflectivity of Cu at this laser wavelength can severely damage the laser sources and optics, leading to high maintenance and operation costs. [START_REF] Jadhav | Influence of selective laser melting process parameters on texture evolution in pure copper[END_REF] Another approach is to print Cu alloys that have much stronger absorption at the laser wavelength than pure Cu. Serval studies report the production of fully dense Cu-alloys such as Cu-Cr, Cu-Sn, or Cu-Zn. [110][111][112] However, a post-processing step is required for the as-built Cu alloys to reach electrical and thermal properties close to those of the bulk Cu. 110,111 Despite these impressive results, SLM of pure Cu, with properties comparable to bulk Cu without post-processing, remains challenging.

This chapter describes the AM of pure Cu using the SLS and SLM technologies. First, a binder-free SLS is developed to fabricate dense, and high-quality Cu parts. For that, a 532 nm Nd:YAG laser is chosen because Cu has a high optical absorptivity (about 40% at 532 nm). Next, the SLM of pure Cu is investigated using an IR fiber laser (λ = 1070 nm) with a moderate power (400 W). Dense, crack-free, and complex shaped heat sinks are fabricated by SLM after optimizing the printing parameters as shown in

Selective laser melting (SLM)

Cu parts were fabricated by SLM using an SLM 125® 3D printer (SLM Solutions). The × 123 × 25 mm 3 ). The density of the samples was measured using the Archimedes method (Mettler Toledo AT201). The microstructure of the printed parts was characterized using an SEM (FEI Quanta 200 Environmental SEM). The cross-sectional views of polished samples were observed using an optical microscope (Nikon Eclipse E600). The average surface roughness was measured using an optical surface profiler (Zygo NewView 8000).

The crystalline phases of the powders and parts printed were analyzed by X-Ray diffraction (XRD, Bruker-AXS D8 Discover Diffractometer, Cu Kα = 0.154 nm). Energy dispersive X-Ray (EDX) was performed on the Cu powders and parts using a Helios NanoLab™ 660 equipped with EDAX TM detector. The electrical resistance of the printed material was measured using a four-point probe method (Keithley 2400 Source Meter) with a current ranging from -0.03 to 0.03 A. The thermal conductivity was calculated from the electrical conductivity using the Wiedemann-Franz law. Sample hardness was measured by Vickers hardness test (Wilson Tukon TM 2500-3), using a load and dwell time of 0.5 Kg and 10 s, respectively. The physical properties of the heat sinks printed were compared to a commercial heat sink (Enzotech Mosfet MOS-C10). The thermal-dissipation performances of the printed heat sinks were evaluated by simulating the thermal management situation of microelectronics using an IR camera (FLIR E85 24º). The heat sinks were attached to an electronic chip with a thermal tape (3M™ 8815 Thermal Tape). Using a hot plate, the memory card and the heat sink were heated to 100, 150, and 200 ºC, respectively, with a fan, which was placed 10 cm above the heat sink for airflow.

Selective laser sintering (SLS) of pure copper

Effect of cold pressed powders in SLS of pure Cu

To enhance the particle contact and thus the surface diffusion, the effect of cold pressing on the SLS process was investigated as illustrate on This phenomenon is attributed to a poor fusing among the loosely packed Cu particles and the splashing of Cu particles during the SLS process. 113,114 The density of the sintered part was estimated to be about 79 %. In contrast, a part fabricated from the cold-pressed pellet has an estimated density of about 91%. The number of pores and the pore sizes decreased significantly. Besides, Cu particles ejected were not observed, and laser scan tracks became continuous in the part fabricated from the cold-pressed pellet. The cold-pressed Cu pellet with a better packing density and more intimate contact between the Cu particles facilitates the SLS process by preventing particle ejection, as shown in The surface roughness of the parts fabricated from the loose-powder bed and the coldpressed was analyzed using a three-dimensional (3D) optical surface profiler. As shown in No change in phase and crystallographic orientations were observed.

Therefore, the average surface roughness of the final Cu parts was strongly affected by the cold-pressing step. Sintering is a diffusion mechanism driven by the reduction of the surface energy. Fick's and Jost's laws describe the diffusion mechanism mathematically and determine its strong dependence on the temperature as well as the material. 115,116 In solid diffusion, different diffusion paths occur at the same time, such as surface, phase boundaries, grain boundaries, and volume diffusion. The faster path is surface diffusion, which depends on the contacts among particles. Next, in order of decreasing speed, is phase boundaries diffusion, which varies depending on the number of phases. Then, grain boundaries diffusion, which is influenced by the grain sizes. Finally, volume diffusion, which is affected by the density of dislocations. [START_REF] Exner | Sintering Processes[END_REF]117 The contact between powder particles was enhanced by the cold pressing step, raising the surface diffusion during SLS and ultimately leading to easier neck formation for densification. Furthermore, it is well known that pressing of powder particles engenders dislocation formation, thus the cold-pressing treatment promoted volume diffusion as well. 118

SLS of dense Cu part using cold-pressed powders

Although a fast diffusion by enhancing both surface and volume diffusions occurred during SLS of a cold-pressed pellet, solid diffusion remains a slow process and therefore not all pores were suppressed. To obtain a fully dense part with a density close to the theoretical limit, the sintering times need to be increased; and thus, the laser scan speed was varied from 125 to 25 µm/s.

Conclusions

The effect of a cold-pressing step on the SLS of Cu powders is investigated. Due to better contacts among powder particles, the surface diffusion is promoted and powder ejection out of the laser tracks is eliminated during the SLS process of a cold-pressed powder pellet.

It is demonstrated that cold-pressed powders, with optimized scan speed, allows fabrication of fully dense Cu parts with a low surface roughness. The whole sintering process is free of binders, which is promising to obtain comparable properties as Cu manufactured by conventional method while enabling complex structure designs without post-printing steps.

Nevertheless, to obtain dense Cu parts, the laser scan speed must be relatively low ( 25μm/s). At a low speed, a long printing time is necessary to manufacture fully dense 3D parts. Thus, this method can only be applied on small objects where low surface roughness is required. For larger objects (i.e., heat sinks), another printing method is required. The next part will present an SLM process in which larger objects can be printed in a shorter time frame.

Selective laser melting (SLM) of pure copper

SLM parameter optimization

Reports on the 3D laser printing of pure Cu show that a high laser power is required to obtain dense Cu parts. The printing parameter windows reported are typically a laser power between 600 to 1000 W (λ = 1070 and 1080 nm, beam size = 40 to 80 μm), a scan speed between 200 to 1000 mm/s, a layer thickness between 0.03 to 0.05 mm, and a hatch distance from 0.05 to 0.12 mm. [START_REF] Colopi | Selective laser melting of pure Cu with a 1 kW single mode fiber laser[END_REF][START_REF] Ikeshoji | Selective Laser Melting of Pure Copper[END_REF][START_REF] Jadhav | Influence of selective laser melting process parameters on texture evolution in pure copper[END_REF] Thus, the printing parameters are optimized in the range of the same parameters as found in the literature to achieve dense Cu parts with smooth surface finishes but by using a moderate laser power of 400 W. The hatch distance is adjusted from 0.08 to 0.12 mm at an increment of 0.01 mm. The scan speed increases from 200 to 600 mm/s with a 50 mm/s increment. The layer thickness is kept at 0.03 mm.

First, the overall shapes of the printed parts is investigated and qualitatively categorized into three groups, which defines the process windows of printing pure Cu. In addition to XRD diffractograms, energy dispersive X-Ray (EDX) was performed on the Cu powders and parts. When the scan speed decreased, sufficient melting of the Cu powders occurred, leading to an adequate flow of molten metal, and thus, a low Ra. 126 The findings obtained by studying the part shape, density, and surface roughness, suggest that at a laser power of 400 W, a scan speed of 400 mm/s, and a hatch distance of 0.12 mm are the optimal parameters to print a crack-free Cu part with a density of 95% and a Ra of 18 μm. To better understand how the printing parameters, influence the density, Ra, and the surface morphology of the printed parts was studied and associated with the molten metal stability. The laser-matter interactions induced the formation of a circular molten pool in several hundred microseconds. 127 Printing parameters, such as the scan speed and hatch distance, influence the dynamic, temperature, shape, and stability of the molten pool as well as the surface finish and the part density. [START_REF] Ikeshoji | Selective Laser Melting of Pure Copper[END_REF]128,129 At a slow laser scanning speed of 200 mm/s, bumps are formed on the surface, resulting in pores formation (Figure 3 Producing high-density Cu parts with a moderate laser power is a critical advancement in additive manufacturing. Moderate powers prevent damage to the laser optics (the back reflection of the laser beam) and lower the energy consumption (60% less power compared to 1 KW laser). 132,133 One possible explanation of the successful additive manufacturing of dense Cu parts with a moderate power is associated with the laser beam. Jadlhav et al.

reported that a low laser power led to porous Cu parts (84%). [START_REF] Jadhav | Influence of selective laser melting process parameters on texture evolution in pure copper[END_REF] Others reported that adding oxide or carbon particles to the Cu parts enhanced the laser absorption, which should have decreased the laser power. Yet, high laser power was still required (> 500 W). [START_REF] Jadhav | Influence of Carbon Nanoparticle Addition (and Impurities) on Selective Laser Melting of Pure Copper[END_REF]134 Several studies demonstrated that the wavelength, beam angle, size, and quality play vital roles in laser-matter interactions. [135][136][137] Naturally, the laser beam characteristics influence the power required to print Cu.

The SLM 125 3D printer is equipped with a 400 W single-mode CW fiber laser (IPG photonics, Yb:YAG, λ = 1070 nm). The laser beam has a stigmatic Gaussian profile, indicating a TEM00 mode, as shown in

3D printing of complex Cu heat sinks

The main advantage of the SLM is its ability to manufacture complex shapes that cannot be achieved by conventional machining. The ability to design thin and sophisticated The optimal printing parameters (scan = 400 mm/s, hatch = 0.12 mm) were applied for the fabrication of the complex structures with smooth, crack-free surfaces. It is worth mentioning that the printing of pure Cu shows a spatial resolution of a few hundred micrometers, which makes the SLM methods a desirable choice for the fabrication of complex components for thermal management.

Mechanical and thermal properties of printed heat sinks vs. a commercial one.

The physical properties of the columnar heat sink printed were measured and compared with the COM columnar heat sink and bulk Cu. Prior to the measurements, the printed and COM heat sinks were sandblasted to obtain a similar surface finish. The densities of the printed and COM heat sinks were 95% and 98%, respectively. Note that the other structures also had densities of approximately 95% (i.e., helix and bent tubes). The resistance of a single Cu pillar was obtained by linear regression of the I/V curve and converted into electrical conductivity using the following equation: 140

σ = L A ×R , ( 6
)
where σ is the electrical conductivity (S/m), L is the length of the pillar (m), A is the area cross-section (m 2 ), and R is the resistance (Ω). The results show that the printed and COM heat sinks have an electrical conductivity of 5.71×10 7 and 5.81×10 7 S/m corresponding to 98 % IACS (International Annealed Copper Standard) and 100 % IACS, respectively. The TC can be obtained from the electrical conductivity measurement for pure metals using the Wiedemann-Franz law: 141

κ σ = KL, ( 7
)
where κ is the TC (W/m.K), K is the temperature (K), and L is a proportionality constant The heat-sink base plate was used for the Vickers hardness tests using a Wilson Tukon TM 2500-3 with a load and dwell time of 0.5 kg and 10 s, respectively. The hardness was calculated using the following equation:

HV ≈ 1.891 × F d 2 , ( 8
)
where F is the force of the intender in N, and d is the average length of the diagonal in mm. Table 1 summarizes the mechanical properties, and electrical and thermal conductivities of the printed and commercial heat sinks and the values of the bulk Cu. It can be seen that the printed heat sink has similar properties to those obtained from the commercial heat sink and bulk Cu. 142,146,147 The high TC of the printed heat sink (i.e., 368 W/m.K) is encouraging. We will explore the heat dissipation of SLM-printed heat sinks that can outperform the COM ones.

Table 1. Physical properties of the printed and the commercial (COM) columnar heat sinks and bulk Cu made by powder metallurgy route.

Heat dissipation performance of the printed heat sinks

The surface area is the main factor in thermal management. SLM has the ability to easily increase the specific surface areas through manufacturing thin features of complex shapes. [START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF][START_REF] Singer | Additively manufactured copper components and composite structures for thermal management applications[END_REF] To demonstrate the benefit of additive manufacturing in thermal management, the heat dissipation performance of different structures printed was measured and compared to the COM heat sink. the highest surface area of 600 mm 2 /g, followed by the helix one with 390 mm 2 /g, and the columnar heat sinks with 286 mm 2 /g. In conclusion, using SLM to print intricately shaped heat sinks yielded dense and crack-free parts with large surface areas and excellent thermal properties. High dissipation performance is achieved with the SLM method and can produce heat-dissipating materials that exceed commercial ones.

Conclusions

In this study, SLM with a moderate laser power (400 W) is successfully implemented to additively manufacturing complex-shaped crack-free functional Cu parts that have a density up to 95% and a surface roughness of 18 μm. A systematic investigation into the printing parameters is performed to optimize the printing quality and understand the part formation mechanism. It is found that the printing parameters play critical roles in controlling the printing quality. Insufficient laser energy deposition leads to rough surface finish associated with bumps and pores, while excess energy deposition leads to destructive balling effects. The high-density achieved in this study is attributed to the high-quality laser beam, which allows to focus the beam with the conservation of the original output power.

Functional Cu heat sinks with various sophistic structural designs are fabricated using the optimized printing parameters. Their physical properties and thermal performance in cooling electronic chips are investigated. These printed Cu parts exhibit a thermal conductivity of 368 W/m.K, which is close to the bulk Cu fabricated using a powder metallurgy process (385 W/m.k).

The thermal performance of the printed Cu heat sinks is compared with commercially ones by studying their cooling dissipation efficiency of a hot electronic chip. The COM and the printed columnar heat sinks with the same dimension show similar cooling performances. However, the heat sinks printed with a helix structure and bent porous tubes outperform the thermal dissipation of the COM columnar structure by 45 % and is attributed to the significantly enhancement of surface areas.

Interphases for Copper Matrix Composite

Introduction

Metal matrix composites (MMCs) need robust interphases to enhance the transfer of properties between the matrix and the reinforcement and thus, tailor the overall properties of the resultant composites. 148 The formation of interphases is particularly crucial in composites comprised of chemically and thermomechanically dissimilar compositions. poor Cu wettability with C due to the lack of chemical affinity. 152 The lack of affinity is more significant when the molten metal is in direct contact with the carbon reinforcement such as in laser-assisted AM. 109,153,154 Extensive experimental and theoretical works have been devoted on the interphase engineering on C-reinforced Cu matrix composite. 155,156 One approach has been to alloy the metal with carbide-forming elements (e.g., titanium (Ti), zirconium (Zr), chromium (Cr), and boron (B)). [157][158][159] However, the properties are severely degraded due to inhomogeneous dispersion of the alloying elements at the interfaces. [160][161][162] The second approach of considerable interest is to coat the carbon reinforcement with carbides. 163,164 Titanium carbide (TiC) offers strong oxidation resistance, high mechanical strength, excellent chemical stability, and low density, making it a right candidate as the interphase. 165 However, it has been reported that TiC does not wet thoroughly with molten Cu (contact angle >115°), resulting in poor Cu-TiC adhesion. 166 Another concern to print Cu/C composites is the oxidation/degradation of the carbon reinforcement. Laser 3D printing involves high temperatures (Tm, Cu = 1085 ºC) and laser matter interactions. Diamonds degrade into graphite at a temperature above 800 ºC, even in a vacuum. Thus, the D particles need to be protected from the molten metal and the laser radiation. Over the past decades, extensive research efforts were undertaken to develop effective oxidation barriers on C materials, such as oxides, nitrides, and carbides. [167][168][169][170][171][172] Among them, refractory carbides have the advantage of spontaneous reaction with C and being permeable to oxygen. 173,174 Substantial evidence suggests that materials coated in multiple carbide layers possess superior resistance to oxidation and thermal ablation as compared to traditional single-layer carbide coatings. 175,176,177 However, the production of multilayer coatings involves complex processes that often require pricey precursors to be decompose at elevated temperatures. [178][179][180] Developing a method straightforward for forming complex multilayer carbide coatings is thus challenging due to the covalent nature bonding of carbides materials, leading to high melting points. 181 In this chapter, a graded Ti oxide-carbide (TiO2-TiC) coating is developed on CFs and D particles to improve the wettability with molten Cu using a molten salt process. In addition, using the same method, an innovative single-step multilayer coating of refractory carbides is presented via an oscillatory reaction (Figure 4.1-1). This process provides a versatile and efficient way to produce transition metal carbides at relatively low temperatures in short time frames. The salt acts as a solvent to dissolve the metal powders and, therefore, enhance the Ti mass transport to the CF surfaces. 182,183 The metal dissolution in the molten salt and mass transport are discussed.

Experimental set-up and procedures

Preparation of coatings

The synthesis of carbide coatings was carried out by a molten-salt process, illustrated schematically in The coating temperature varied from 800 to 950 ℃ for times of up to 5 h. These coating conditions were chosen based on the study by Lui et al.. 184 After cooling down, the crucible was immersed in a boiling distilled water (~100 ºC) until the salt was fully dissolved.

Finally, the coated C were filtered from the solution, flushed with deionized (DI) water, and air-dried in an oven at 60 ºC for 1 h. The filtration dissolution/filtering process was repeated three times.

Wettability analysis

The wettability of molten Cu on the coated CFs was investigated using an approximate equal-molar mixture of 0.2 g (~0.017 mol) of coated CFs with 1.2 g (0.019 mol) of dendritic Cu powder (35 µm, ECKA Granules Germany CmbH). The mixture was placed in an Al2O3 crucible with a carbon paper liner to prevent reaction between the Cu powders and the crucible. The crucible was put into a tube oven filled with flowing argon gas and heated to 1200 ºC with a rate of 10 ºC/min. The crucible was held for 10 min and then cooled to room temperature at a rate of 15 ºC/min. Contact angles between the CFs and the Cu droplets were determined using scanning electron microscopy (SEM) and cross-sectional images of resin-encased and polished CFs.

Quenching tests

Rapid thermal quenching of the molten salt in the ambient air was used to terminate the coating process at intermediate stages. This enabled the study of the carbide formation, including metal dissolution, diffusion, and reaction with C. The quenching tests were performed by heating the powder mixture to 800 and 950 ºC for 0.5 and 5 h, respectively.

The hot crucible was then rapidly quenched in the air at a cooling rate of about 500 ºC/min, resulting in KCl (mp = 770 ºC) solidification in less than 20 s. A piece of the solidified KCl-CFs-Ti and KCl-CFs-Cr-Ti samples were embedded in a polyphenolic resin, and mechanically dry polished for SEM analyses.

The growth rates of the TiC and Cr3C2 layers

The diffusivity and associated activation energy were measured by mixing KCl, CFs, and Ti or Cr with molar ratio of KCl:CFs:Ti = 0.642:1:0.238 and KCl:CFs:Cr = 0.642:1:0.238, respectively. Both coating times and temperatures were varied between 1 to 5 h at processing temperatures of 800, 850, 900, and 950 ºC, respectively. After the reactions, the coated CFs were washed and dried as described earlier, before embedded in a polyphenolic resin and polished for SEM analyses.

Erosion/oxidation under an oxy-acetylene flame

A combustion torch with a 1.5 mm orifice was used to produce a combustion flame. The precursor gases were a mixture of 99.6% acetylene (C2H2) and 99.9% oxygen (O2) with flow rates of 1780 and 1800 sccm (standard cubic centimeters per minute), respectively.

Coated CFs were glued onto a tungsten carbide (WC) substrate containing 6% of cobalt 4.3 Graded titanium oxide -carbide coatings on CFs.

Effect of temperature on TiCx coating composition and morphology

The effects of the reaction temperature on the coating morphology and structure are shown in In contrast, coatings at 900 and 950 ºC have far fewer and smaller pores and appear almost fully consolidated. In these cases, the temperature is sufficiently high that localized differences in surface reactivity become less critical for determining the growth morphology. Moreover, the possibility of sintering driven by surface diffusion would be more significant. Both effects, individually or combined, could explain the dense coatings obtained at high temperatures. Others have also reported consolidated coatings at temperatures above 900 ºC in different salts. 164,187 In addition to the surface morphology, XRD was performed at each processing temperature to quantify the product phase composition ( The coating porosity was estimated from a mass balance of the reactants as follows:

𝑥𝑚𝑜𝑙𝑒 𝐶 𝑟𝑒𝑎𝑐𝑡𝑒𝑑 = 𝑚𝑜𝑙𝑒 𝑇𝑖 𝑟𝑒𝑎𝑐𝑡𝑒𝑑 = 𝑚𝑜𝑙𝑒 𝑇𝑖𝐶 𝑓𝑜𝑟𝑚𝑒𝑑 , (9)
Thus, for a CF coated with TiCx, the mole can be expressed as:

𝑥2𝜋𝐿 𝐹 (𝑅 0 2 -𝑅 𝑐 2 )𝜌 𝑐 𝑀 𝑐 = 2𝜋𝐿 𝐹 (𝑅 𝑇𝑖𝐶 2 -𝑅 𝑐 2 )𝜌 𝑐 𝑀 𝑇𝑖𝑐 (𝜌 𝑇𝑖𝐶 (1 -𝜑)), (10)
where x is the mole fraction of C in TiCx product (0.93), LF is the length of the fiber ( around 200 µm), MTiCx and MC are the molecular weight of TiCx and C respectively (59.9

and 12 mol/g), ρTiCx and ρC are C and TiC densities (1.68 and 4.93 g/cm 3 ), and φ is the coating porosity. The outer radius of the coated CF, the radius of the starting CF, and the CF core radius at the end of the reaction are RTiCx, R0, and RC, respectively. After rearranging and combining the terms:

𝑅 𝑇𝑖𝐶 = [𝑘 ′ (𝑅 0 2 -𝑅 𝐶 2 ) + 𝑅 𝐶 2 ] 1 2 , (11)
where

𝑘 ′ = 𝑥𝑀 𝑇𝑖𝐶 𝑀 𝐶 𝜌 𝐶 𝜌 𝑇𝑖𝐶 ×(1-𝜑) , (12)
Similarly, the radius of the unreacted CF is given by:

𝑅 𝐶 = 𝑅 0 (1 -𝑥𝑍 𝑇𝑖/𝐶 ) 1/2 , ( 13
)
where ZTi/C is the mole ratio of Ti to C in the starting reactants. The reactant mole ratio (ZTi/C) used to prepare the coated CFs shown in ). [188][189][190] The results show, within the experimental errors, that the stoichiometry is ~TiC0.73 for the product formed between 800 to 900 ºC. On the other hand, the coating produced at 950 ºC has a higher C content with a stoichiometry of ~TiC0.93.

One explanation for the change is that the Ti and C diffusivities were both greater at higher temperatures and thus tended to form a more reacted coating. This explanation is supported in part by the measurements of the self-diffusion coefficients of C and Ti in TiC as a function of the temperature, [191][192][193] as well as results from the process kinetics, described later in this chapter.

Interphase elemental compositions (TiCx vs. TiOy-TiCx)

In the previous section, it was shown that the salt-bath temperature influenced the morphology and stoichiometry of the TiCx coatings. This section will describe how a combination of SEM, AES, and XPS were used to quantify the chemical compositions of the coated layers produced at the low (800 ºC) and high (950 ºC) temperatures. The AES compositional mapping in The spectra of C1s show a higher intensity of C-Ti for CFs with TiCx than for CFs with TiOy-TiCx, suggesting a higher concentration of TiC on the surface where the same tendencies can be observed on the Ti 2p spectra. Moreover, it can be noticed that the intensity of the Ti-O peaks decreases when the coating temperature increases, suggesting a higher amount of Ti-O at a lower temperature accompanied by a lower concentration of Ti-C. Finally, on the O1s spectra, a higher intensity of O-Ti is observed for the CFs with TiOy-TiCx. Standard spectral analysis procedures were used to deconvolute the XPS Ti 2p spectra into contributions from TiCx and TiOy (Table 4.3-1). The results display the same trend as the AES results: the distributions of surface Ti bonded as Ti-C vs. Ti-O strongly depend on the coating temperature. The results demonstrate that ~50% of the surface Ti is bonded as TiOy at 800 ºC but only ~10% at 950 ºC.

Table 4.3-1 XPS quantitative analyses using peak integration.

A possible explanation is that the TiOy formed during the boiling-water (~100 ºC) dissolution of the KCl salt at the end of the process. Naturally, the effective oxidation depth (via both the AES and XPS) would be larger for the highly porous coatings obtained at 800 ºC, simply due to the enhanced water penetration. Also, the coatings obtained at 800 ºC have a more stoichiometrically Ti-rich layer (TiC0.73, as described in the previous section), suggesting that more reactive, less -bonded Ti sites could be oxidized. In contrast, the highly consolidated and essentially pure TiCx coatings formed at 950 ºC (i.e., TiC0.93)

would allow much less water penetration and, thus, less oxidation.

Cu wettability and interphase surface energy

To understand the role of a graded TiOy-TiCx coating to tailor the interfacial bonding between Cu and C, we prepared two types of CFs coated with either a high or a low concentration of oxide content, respectively. The contact angles (wettability) of Cu on coated CFs prepared at 800 and 950 ºC were measured. The well-known Young-Dupre equation was used to estimate the adhesion energy:

𝐸 𝑆𝐿 = 𝛾 𝐿𝑉 (1 + 𝑐𝑜𝑠 𝜃), ( 14
)
where Esl is the energy (work) of adhesion (mN.m -1 ), lv is the liquid/vapor surface energy (tension) (mN.m -1 ), which has been reported to be 𝛾 𝐶𝑢 = 1330 -0.

where ρTi is the density (4.59 g/cm 3 ), MTi is the molar weight (47.9 g/mol), R0Ti is the radius of Ti particles (20.10 -6 cm), DTi n+ is the diffusion coefficient for Ti n+ in molten KCl.

The diffusion coefficient can be estimated from the viscosity () by the well-known

Stokes-Einstein equation:

D ~ kt 6πηr h , ( 16
)
where rh is the effective hydrodynamic radius of the Ti n+ and k is Boltzmann's constant.

The reported viscosity of molten KCl at 800 ºC is quite low (~1.2 mPa-s). 200 Thus, assuming an effective Ti n+ radius of ~0.1 nm that yields a diffusion coefficient of 8.10 -6 cm 2 /s. Finally, ΔCTi = Csat -Cb where Csat is the concentration of Ti n+ at saturation and Cb in the liquid KCl. We are interested in estimating the time to dissolve a Ti particle in the early stage of the process. Thus Cb = 0 and the following simplification can be made:

∆C Ti = C Sat , (17)
Sekimoto et al. reported the saturation concentration of Ti 2+ and Ti 3+ in KCl-NaCl at 740 ºC are of about 1.10 -3 to 1 mol/l. 197 By using the lowest Ti n+ concentration, the result shows the dissolution time is concise and of about 4 min.

The rate of diffusion mass transport of Ti n+ to the outer CFs or TiCx-CF interphase surface (step 2) is controlled by the viscosity of the KCl and the mean distance between the Ti o particles and the CFs. Assuming a statistical average distribution of particles, the mean interparticle distance is, conservatively, smaller than 100 µm. The approximate time for Ti n+ to diffuse this distance can be estimated from:

t d ~<x> 2 2D Ti n+ , (18)
Eq (18) gives an estimated time of td < 10 s to diffuse the 100 m (mean Ti 0 -CF interparticle distance). The conclusion from the simple analysis above is that both the particle dissolution (step 1) and Ti n+ diffusion mass transport (step 2) are fast. Thus, the process step governing the rate of TiCx coatings formation is step 3, the diffusion and mass According to classical kinetic theory, the coating thickness can be described as a function of the square root of the coating time using the following equation: 201,202 x = √D eff T ,

where x is the carbide coating thickness (cm), t is the coating time (s), and Deff is the effective diffusivity through the TiCx layer (cm 2 /s). As shown in The effective diffusivity is expected to show an Arrhenius temperature dependence by an activation energy: 204,205

D eff = D 0 e -Q eff RT , ( 20
)
where D0 is the pre-exponential factor, Qeff is the activation energy (J/mol) governing the process and R is the gas constant (J/mol.K). The effective process activation energy was determined to be ~88 KJ/mol by linear regression to the data in Here the Kirkendall effect related to the rapid counter diffusion of C that ultimately generates voids at unreacted carbon interfaces. 208 The diffusion rates of C in TiC have been reported to be several order of magnitude faster than Ti in bulk TiC. 193 Consequently, the more rapid diffusion of C than that of Ti would be expected to generate Kirkendall like voids at the reaction interface. Moreover, higher void generation was observed at 950 ºC than 800 ºC.

This suggests that the C diffusion rate increases with the temperature. 191,193 EBSD analysis was performed on the TiCx coating obtained at 950 ºC (Figure 4.3-14 (f)) and showed a polycrystalline columnar structure with an average grain size close to 1 µm. The grains arrange in a conformal structure. The grain boundaries are aligned along the radial direction. These grain boundaries may serve as an easy diffusion path for C and thus enhance the C diffusion rate at higher temperatures.

Graded TiO2 -TiC coatings on diamond particles

In the previous part, we studied the impact of molten-salt-bath temperature on the TiC coatings on CFs. It was demonstrated that a low bath temperature at 800 ºC resulted in the formation of a TiO2-TiC layer. Compared to a TiC coating on CFs, the TiO2-TiC surfaces facilitated the adhesion of Cu particles on CFs. In this part, we performed a similar study to reveal the influence of molten-salt-bath temperature on the coatings on diamonds.

Conclusions

In summary, tailored TiOy-TiCx coatings are synthesized on CFs. The TiOy formation is observed after the dissolution of the molten salt in boiling water and is assumed to be the The Ti-based coating will be used later in the laser 3D printing of Cu/D composite materials.

Multilayer coatings of refractory carbide in a molten salt

Single-step coating of multilayer carbides in a molten salt

Creating a multilayer carbide coating could efficiently and effectively protect carbon materials under extreme environment (i.e., laser 3D printing). For that, the multilayer coating process necessitates a salt media which can dissolve several metallic powders; in our case Cr and Ti powders. The literature shows that a eutectic salt composed of sodium chloride and potassium chloride (NaCl:KCl) is suitable to coat carbon materials with Cr3C2

while KCl allows the formation of TiC layers on the CFs. 184,211 Therefore, KCl appears to be an excellent candidate as a reaction media to synthesize both TiC and Cr3C2 layers on CFs. To confirm the formation of these carbides, the CFs were coated using KCl at a time and temperature of 5 h and 950 ºC, respectively. After coating and washing away the salt, the CFs were analyzed by SEM using the back-scattered electron (BSE) mode to enhance the contrast between the CFs and the layers coated (Cr3C2 appears to be brighter than TiC which is brighter than C). 212

Oscillatory reaction in a molten salt

The liquid salts containing a mixture of Ti + Cr and CFs were frozen at different coating steps using an air quenched method to understand how the layer sequence is formed in the molten KCl.

The kinetic analysis of the multilayer coating

To clarify the formation of the first Cr3C2 layer during the early stage of the coating, the kinetic growth of single carbides (M= Ti, M = Cr) was studied independently at different temperatures and times ranging from 800 to 950 ºC and 1 to 5 h, respectively, as shown in Therefore, the single-layer coating of TiC or Cr3C2 is not suitable as oxidation barrier coating on CFs. Delamination and pores were observed on coatings resulting in the partial or total oxidation of the CFs. Nevertheless, when TiC and Cr3C2 were combined as a multilayer coating, an intimate CF-coating interface is observed. In addition, the fiber's graphitization was estimated based on the D-band (~1350 cm -1 ; related to defects and disorder in the carbon structure) and G-band (~1580 cm -1 , results from ordered graphitic structure) ratio (ID/IG) (Figure 4.4-8). 221 It is found that ID/IG is of about ~ 0.90 for the original and after the flame exposure CFs. Hence, the high graphite crystallinity was conserved and demonstrates the effectiveness of the multilayer coatings as a protection barrier for carbon materials.

Conclusions

In summary, a molten-salt coating process is investigated using Ti, Cr, and Ti + Cr to produce homogenous coatings of single and multilayer carbides on CFs at a relatively low temperature (950 ºC). The microstructural characterization of the complex multilayer carbides shows a three-layered coating composed, starting from the CFs, of an inner Cr3C2 layer, followed by a TiC layer, and an outer Cr3C2 layer. The study confirms the advantage of the molten-salt process as a single-step, cost-effective coating method for synthesizing multilayer carbides on CFs.

Furthermore, analyses of the coating process show an oscillatory reaction mechanism of forming the multilayer carbide coating. In the early stage of the process, the liquid salt is saturated with Ti and Cr, which react with the CFs to form carbide layers. Subsequently, Ti began reacting with Cr particles to form a shell, which prevents further Cr dissolution.

Once a complete TiC layer formed, Cr is free to dissolve inside the liquid salt and react with the CFs to create the final Cr3C2 layer. Additionally, the kinetic growth of TiC and Cr3C2 show that the Cr3C2 layer grew twice as fast as the TiC layer, implying a faster reaction/diffusion of C species in Cr. Thus, the inner Cr3C2 layer resulted from a kinetic regime at a high temperature (950 ºC).

Finally, the advantages of the multilayer over a single carbide coating are demonstrated by high-temperature oxidation under an oxy-acetylene torch. The multilayer coatings composed of TiC and Cr3C2 combined the physical properties of both materials to create an efficient and effective barrier layer for CFs in harsh environments.

Additive Manufacturing of Cu / D Composite Materials

Introduction

The manufacturing of materials with high thermal conductivity (TC) and intricate shapes is crucial for the development of highly efficient heat-exchanging components for microelectronic industries. [START_REF] Vasiliev | Heat pipes in modern heat exchangers[END_REF][START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF] Creating a composite material consisting of two highly thermal conductive elements, such as Copper (Cu) and Diamond (D), may fulfill these unmet needs. [START_REF] Qu | Review of metal matrix composites with high thermal conductivity for thermal management applications[END_REF] However, the lack of chemical affinity between Cu and D and the difficulty of machining D-based materials presents significant challenges. [START_REF] López | The solubility of C in solid Cu[END_REF][START_REF] Silvain | Copper-Carbon and Aluminum-Carbon Composites Fabricated by Powder Metallurgy Processes[END_REF] Extensive research has been devoted to achieving highly thermally conductive metal matrix composites (MMCs) by promoting the interfacial bonding between Cu and D. 222,223 Common interphases are composed of carbide elements and have demonstrated enhanced thermal properties compare to pure Cu. 163,[224][225][226] However, conventional manufacturing methods, such as high-pressure sintering, spark plasma sintering, or infiltration, suffer several limitations when sophisticated geometries are required.

The recent development of laser 3D printing has attracted much attention due to the ability to design complex shapes and has exhibited impressive results on metals and alloys. [START_REF] Tran | 3D Printing of Highly Pure Copper[END_REF][START_REF] Frazier | Metal Additive Manufacturing: A Review[END_REF]109,111,112 Over the past decades, different approaches were undertaken to laser 3D print MMCs via selective laser melting (SLM). The first approach lies in simply mixing the reinforcement and metal powders. However, the large size of the reinforcement modifies the molten pool dynamic and rheology, leading to the fabrication of porous parts. [227][228][229] The second approach involves the creation of a composite powder containing the metal matrix and the reinforcement to form a homogenous mixture at the microscale. [230][231][232][233] This method allows to 3D print dense MMCs with mechanical properties comparable to ones obtained via conventional method. Yet, only nanocomposites can be manufactured and, thus, cannot be generalized for micro-reinforced MMCs. The third approach consists in the in-situ formation of the reinforcement. [234][235][236][237] This technic enables obtaining dense MMCs with micro-size reinforcement, but a limited number of reinforcements can be formed during SLM. Although numerous attempts to laser 3D print MMCs, to our best knowledge, additive manufacturing (AM) of Cu/D composites remains less explored.

In this chapter, the feasibility of laser 3D printing of Cu/D composite materials is investigated via laser directed energy deposition (LDED) as illustrated in Figure 5.1-1 (a).

The printing parameters are optimized to obtain a dense material. Also, the TC of the printed composite as well as the quality of the D after the process are studied. Later, the fabrication of complex Cu/D composite 3D structures is studied via SLM (Figure 5.1-1 (c)). It is shown that the presence of D affects the molten pool dynamics. Different printing strategies are employed to overcome this issue. Finally, highly complex designs are printed which have potential used in thermal management applications.

Experimental set-up and procedures

Laser directed energy deposition (LDED)

D powders with a mean diameter of 105 μm (Eastwind Diamond Abrasives) were coated with a graded TiO2-TiC layer using a molten-salt process. The energy density (J/mm 3 ) was calculated with the following equation:

𝐸 = 𝑃 𝐿𝑇×ℎ×𝑠 ( 21
)
Where P is the laser power (W), LT is the layer thickness (mm), h is the hatch distance (mm), and s is the scan speed (mm/s). The crystalline phases and chemical structures of the coated D were analyzed both by X-ray diffraction and micro-Raman spectroscopy. The volume percentage of D in the printed composites was estimated by image analyses using

ImageJ. The density of the printed samples was measured using the Archimedes method.

The microstructures of the parts printed were characterized using a scanning electron The method consists of heating the front face of the sample with a pulsed laser and to measure the related temperature increase on the opposite face using a fast (10 MHz) IR detector. The heat transfer is thus 1D with the spatial resolution is the bulk material. The material isotropy was considered. The measurement has been performed along with the two directions of the sample and no difference was noticed, applying that the building direction does not affect the thermal conductivity. Comparable TC was measured with a 2% deviation range. Also, a photodiode was set to trigger the acquisition after each laser pulse. An averaging process was performed over the period of repetition that allows to increase the signal-to-noise ratio of the response by the detector. The TD of the sample is identified by minimizing the gap between the experimental response and the theoretical one calculated using an analytical solution of the heat transfer within the material. The minimization is achieved by implementing a nonlinear least square algorithm (the Levenberg-Marquardt technique). The TD was converted into TC using the following equation:

k(T) = a(T) × C p (T) × ρ(T), ( 22
)
where k is the TC (W/m.K), a is the TD (mm 2 /s), ρ is the density (g/cm 3 ), and Cp is the heat capacity of the sample (J/Kg.K). The Cp was calculater using the rule of mixture using the following value : D = 630 J/Kg.K, TiC = 190 J/Kg.K and Cu = 392 J/Kg.K. 239

Selective laser melting (SLM)

Coated-D particles (Eastwind Diamond Abrasives, 25 µm) were mixed with spherical metallic Cu powders (US Nano, Cu 99.5%, mean diameter 45 µm) at a volume percentage of 5% (Figure 5.1-1 (d)). The SLM was carried out using the SLM®125 3D printer from SLM Solutions as described in Chapter 3. A stipe laser scan strategy was used for all prints.

The laser power and layer thickness were kept constant throughout the study to 400 W and 0.03 mm, respectively. First, a single melting strategy was investigated for a scan speed and hatch ranging from 100 to 600 mm/s and 0.03 to 0.2 mm, respectively. The single melting consists of shining the laser once on each layer, as illustrated in

Laser directed energy deposition of Cu/D composite materials

Optimization of the printing parameters

The realization of dense Cu/D composites with high TC requires the formation of an interphase between the Cu and D because of their chemical incompatibility. [START_REF] Silvain | Copper-Carbon and Aluminum-Carbon Composites Fabricated by Powder Metallurgy Processes[END_REF]211 In Chapter

Interfacial characterization of Cu and coated-D

The interfacial bonding between the matrix and the reinforcement is one of the main concerns to obtain dense and high-TC composites.

Diamond quality through the printing process

The D quality is a crucial point to study the feasibility to laser 3D printing of Cu/D composites due to the potential D graphitization during the printing process. The D quality was evaluated after being immersed in a salt bath (coating process) and exposed to a highpower laser (printing process) as shown in The as-received D powder showed a sharp peak around 1331.5 cm -1 , with a full width at half maximum (FWHM) of 1.9 cm -1 , confirming its high degree of crystallinity (Figure

5.3-3 (d) and (e))

. 245,246 D particles were recollected and reevaluated after the salt coating and printing processes. As shown in

Conclusion

This work demonstrates the feasibility of additive manufacturing using Cu/D

composites. An energy density of about 1250 J/mm 3 is necessary to make dense composite materials with a limited amount of printing defects and high TC (330 W/m.K). The introduction of the Ti-based coating on the D particles before printing enables the formation of a chemical bonding between the matrix Cu and reinforcement D, facilitating the printing of dense Cu/D composites. In addition, the quality of the post-printing D is found to be close to the original one, making the process suitable for manufacturing high-TC materials. However, LDED leads to the formation of coarse cubes and may not be suitable to fabricate intricated heat sink designs. Recall that, in Chapter 3, a dense Cu parts were printed via SLM with a laser power (P), layer thickness (LT), scan speed (s), and hatch distance (h) of 400 W, 0.03 mm, 400 mm/s, and 0.12 mm, respectively. Also, Y. Ma et al. studied the SLM of Al12Si/D composites and found that the energy density necessaries to manufacture MMCs must be higher than that for the puremetal. 228 Therefore, to optimize the print, the following parameters were investigated on Cu and D powders: printing strategy: single melting; s = 100 mm/s to 600 mm/s; and h = 0.03 to 0.20 mm. The laser power and layer thickness are fixed to 400 W and 0.03 mm, respectively, for all experiments. enhances the molten metal flow and the wettability with the previously melted layers, resulting in a denser materials. 248,249 In addition, the energy required to 3D print Cu/D composites appears more substantial than that for pure Cu, where an energy density of 275 J/mm 3 led to a dense material (95%). One explanation is that the presence of solid coated-D particles into the molten Cu increases its viscosity and limits its ability to flow and fuse. 228,231 Also, Cu is known to have a high optical reflection in the near-infrared domain (i.e., λ = 1070 nm). [START_REF] Hess | Continuous Wave Laser Welding of Copper with Combined Beams at Wavelengths of 1030nm and of 515nm[END_REF]109 When setting an energy density above 1000 J/mm 3 (i.e., low scan speed < 200 mm/s), important back-refection into the laser occurred (>10 W), which triggered the laser back-reflection security and make it shut down automatically.

Selective laser melting (SLM) of

Consequently, a lower scan speed limit of 200 mm/s must be fixed to prevent damage to the laser optics and to the laser itself. Therefore, the maximum density which can be reached using a single melting is about 88%. These observations imply that the coated-D particles have a major effect on the pool dynamics. Two types of spatter particles are categorized in SLM. [START_REF] Matthews | Denudation of metal powder layers in laser powder bed fusion processes[END_REF] During the laser-matter interaction, a spheroidal molten pool is created within a few microseconds. 127 Quickly, a metal vapor appears above the molten pool, inducing a recoil pressure on the pool surface.

Meanwhile, inside the liquid, a thermal gradient rapidly emerges, which inducing a flow of the molten metal from the hot area to the coldest one, known as the Marangoni effect.

Both effects represent the cause of particle ejection, called recoil-induced ejection, as illustrated in Figure 5.4-3. 129,250 In the meantime, a denudation zone is created on the sides of the molten track where powders can be swept under the flow of the metal vapor and the gas shield, named entrainment particles. In this scenario, the particles can be "cold" if there is no interaction with the laser beam or "hot" if they intercept with it (Figure 5.4-3). [START_REF] Ly | Metal vapor micro-jet controls material redistribution in laser powder bed fusion additive manufacturing[END_REF] Naturally, the presence of solid coated-D particles in the pool changes the molten pool rheology, surface tension, and the metal vapor density, which explain the important particle ejections as compared to a pure Cu powder bed. Furthermore, spatter particles affect the quality of the printed parts. These particles fall back into the powder bed after cooling down. Later, during the deposition of a subsequent layer, they partially fuse with the solid material which leads to the formation of bumps, pores, and cavities. Moreover, if spatter particles or the part roughness is higher than the layer thickness, damage of the recoater blade occurs, and uneven powder distribution arises. [START_REF] Wang | Mechanisms and characteristics of spatter generation in SLM processing and its effect on the properties[END_REF] Several studies showed that the laser power and the scan speed are major parameters that affect the number and size of spatters at constant layer thickness, pressure and gas flow. 129,251 It is commonly accepted that a faster laser scan speed lowers the molten pool temperature, vapor pressure and decreases the denudation zone, thus limits the spatter ejection. Nevertheless, as shown before, a rapid scan speed leads to an insufficient flowing and fusing of the powders; thus, porous materials are formed. Hence, the conventional SLM (i.e., single melting) of Cu/D composites raises several issues. First, the high back reflection of Cu in the near-IR domain limits the laser scan speed to a minimum of 200 mm/s. On the other hand, the presence of coated-D particles in the powder bed perturbs the molten Cu dynamics, generating spatter ejections and leads to the porous parts during SLM.

Powder-bed quality for different printing strategies

The successful manufacturing of a micro-reinforced MMCs requires to limit or suppress the spatter ejections. For that, two different printing strategies were introduced in the SLM and their effect on the spatter ejection were monitored and compared with the single melting situation using the layer control system (LCS). Note that the LCS images allow observing only significant particles that represent a small amount of the spatter ejection.

Only large particles are considered here, as they are most likely to influence the printing quality. (i.e., molten pool) are often a few times larger than the raw powders, while entrainment particles are about the size of the bed powders. [START_REF] Ly | Metal vapor micro-jet controls material redistribution in laser powder bed fusion additive manufacturing[END_REF] Recall that the metal vapor was found to be more critical for a Cu/D powder bed than the Cu one for the same printing parameters.

Hence, it is reasonable to assume that these particles originate from an instability of the molten pool due to an increase in the recoil pressure.

Consequently, a recoating strategy was implanted during the SLM to improve the density and part quality, which consists of adding more powder to the current layer by repassing the recoater and shinning the laser for the second time (

Quality and density of printed composite with different printing strategies

Study of the composite qualities associated with different printing strategies was carried out for scan speeds and hatch distances ranging from 200 to 450 mm/s and 0.06 to 0.12 mm, respectively. Photographs of the printed samples can be seen in

Microstructures of Cu/D composites printed with different strategies

The influence of spatter on the layer-upon-layer manufacturing of MMCs was then correlated to the sample densities by observing their cross-sectional and surface Note that the section was polished and slightly etched to reveal particle inclusion inside the composite. The sample printed with a unique fusion shows pores, cavities, partially fuse inclusions, and "balls". The presence of these printing defects was expected, considering the previous observations. Undoubtedly, the addition of new layers on highly defected Cu/D layers leads to uneven powder dispersion. Hence the defects remained during the printing, contributing to the formation of the low-density materials. As mentioned previously, the recoating approach allows filling the pores and cavities during the SLM by adding a new powder layer and shining the laser the second time.

Nevertheless, spatter ejection continues, adding up defects to the composite material. Thus, a partially dense sample was obtained, as shown on the SEM micrographs. 6 Diamond coating by chemical vapor deposition assisted ultraviolet lasers

Introduction

Diamond (D) has been referred as "the ultimate engineering material" for a wide range of applications which are attributed to its exceptional thermal, optical, mechanical, and electronic properties. 262 Chemical vapor deposition (CVD) of diamonds at low pressure has been extensively studied over the past decades due to its numerous advantages, including low cost, high growth rate, and high diamond quality. 263,264 All CVD methods, including plasma, hot-filament, and combustion-flame CVD, involve a number of chemical reactions of hydrocarbon precursors near thermal equilibrium, [265][266][267][268] where a lack of selectivity among different reaction channels take place, towards products or desired products.

With the significant advances in laser technology, lasers provide a unique means for efficient dissociation of precursors or radicals by exciting specific transitions (electronic, vibrational, rotational) in reactant molecules, enabling reaction control to some extent. 269- 271 Laser chemistry has long been explored as a means of seeking chemical control of molecular reactions.

Diamond film characterizations

Surface morphologies and cross-sectional microstructures of the diamond films were characterized using SEM. The growth rate was calculated by dividing the film thickness by the deposition time. D quality was evaluated using a micro-Raman spectrometer. TEM was performed to study the growth transition zone at the film-substrate interface.

6.3 UV laser-assisted (KrF, λ = 248 nm) combustion synthesis of diamond films

Effects of UV laser irradiation on the flame chemistry

Material synthesis arises from chemical processes occurring in the gas phase. 279 Examination of the gas phase variation under UV irradiation is, therefore, critical to determine its role in the growth process. As one of the most-used techniques for detecting excited species in a reaction system, OES of the diamond-forming flames under different laser irradiation conditions was studied.

v'=1) → X '3 P (v"=1)
transition, detection of photons is allowed with a wavelength of 512.9 nm, as indicated in Obvious increments in the concentrations of both excited and ground-state C2, CH, and OH species were observed in the excimer-laser-irradiated flames. These species are all considered to be critical for diamond formation. 266,[282][283][284][285] To determine how these species were generated, the flame temperature, Tflame, was analyzed at different laser fluences. Due to a strong coupling between the translational and rotational energy states, the flame temperature can be approximated by the rotational temperature derived from the emission intensities of the high-resolution rotational lines of CH using the Boltzmann plot (Figure

6.3-5 (a)

). 286,287 The Boltzmann plots were derived from the spectra using the following equation:

ln ( Iλ 4 S J ′ J ′′ ) = - 1 T f E J ′ k + lnC, ( 23
)
where I is the relative emission intensity of a rotational line obtained from the experimental spectrum; C is a proportionality constant that is the same for all rotational transitions within a band; SJ'J" is the rotational intensity factor, l is the wavelength of the emitted spectral line, EJ' is the rotational energy of the initial level, k is the Boltzmann constant, and Tr is the rotational temperature. molecules by either 193 or 248 nm laser have also been reported. 273,275 The higher the laser fluence is, the larger the photon flux would be, leading to the generation of more reactive species. The electronically excited reactive species are generated from the direct UV photolysis in a "nonthermal" fashion.

Effect of KrF (λ = 248 nm) excimer laser irradiation on diamond growth rate and quality

Based on the understanding of the flame gas phase variation with the KrF excimer laser irradiation, we prepared diamond films under KrF excimer laser irradiation at a fixed laser frequency of 35 Hz with laser fluences tuned from 0.6 to 1.4 J/cm 2 to study the influence of the KrF excimer laser irradiation on the diamond growth. The surface morphologies of the diamond films deposited at different laser fluences, from 0.6 to 1.4 J/cm 2 , are shown in Such a narrow transition zone was comparable to the D films on WC prepared using microwave-enhanced and hot-filament CVD, as reported, 295,296 which are known to produce better quality D films than combustion CVD. The narrow transition zone observed with KrF excimer laser irradiation confirms again that KrF excimer laser irradiation improves the diamond growth process by suppressing the non-D carbon formation. The observation of fast lateral grain size involution and the significantly shortened nucleation time can be explained by the observation of gas phase variation in the flame.

We found a pronounced increase of OH-abundance in the center region of the flame. OHradical plays a critical role in combustion synthesis of D by etching the surface-bond hydrogen and stabilizing the sp 3 hybridized surface C bonds. 273,274 Its increment explains why the graphitic C accumulation was greatly suppressed with KrF excimer laser irradiation. Furthermore, another possible contribution to the significantly promoted diamond growth is the increased atomic H concentration due to promoted H-abstraction processes with KrF excimer laser photolysis. The photo-induced dissociation of hydrocarbons preferentially occurs at C-H bonds through H-abstraction processes rather than at C-C bonds. 297 The primary processes of 193 nm UV photolysis of C2H2 are:

C 2 H 2 → C 2 H + H, (24)
C 2 H 2 → C 2 + H 2 . 298 (25)
The primary processes of 147 nm UV photolysis of C2H4 are:

C 2 H 4 → C 2 H 2 + H 2 , (26)
C 2 H 4 → C 2 H 2 + 2H. 299 (27)
The reactions of atomic H not only control the gas-phase chemistry but also determine the availability of reactive sites upon the sites. D growth from oxyacetylene flame proceeds through two critical surface reactions: 1. addition of reactive hydrocarbon radicals on the active surface sites; 2. H-abstraction from the hydrocarbon radicals to make more reactive sites to accept hydrocarbons and stable sp 3 hybridized carbon bonds. 289 The promoted Habstraction processes with KrF excimer laser photolysis of hydrocarbons thus could play a dominant factor in promoting diamond growth. However, more detailed study is required to confirm this hypothesis because the KrF excimer laser used in this study has a wavelength of 248 nm, longer than those used in the reported photochemistry study of C2H2 and C2H4, making the photo-induced H abstraction rate much smaller. 298,299

Conclusions

The influence of UV-laser-induced photolysis in D growth is investigated. OES and LIF of the flame indicated that the UV-laser-induced photolysis produced large quantities of reactive radicals in the flame, contributing directly to the promotion of D growth.

Investigation of the nucleation process suggests that the KrF excimer laser irradiation modifies the diamond growth in a way favoring diamond formation and suppressing nondiamond C accumulation, leading to an enhanced diamond deposition rate and improved diamond quality. The transition layer is reduced from 700 nm without laser to 5 nm with UV laser irradiation. The discovery of the laser photochemistry advantages in diamond growth is of great significance for extensive improvement in the synthesis of a broad range of technically important materials. The D quality was evaluated from the Raman spectra by calculating the quality factor. 292 It is worthy to note that the laser fluence has a linear influence on the growth rate and D quality at 248 nm, but a super-linear influence at 193 nm. Moreover, a larger grain size, a higher deposition rate, and a higher D quality factor were obtained with UV laser irradiation at 193 nm than those obtained at 248 nm, although the laser fluence of the 193 nm laser was one order of magnitude lower than that of the 248 nm laser. The results suggest the advantage of a short laser wavelength from the irradiating laser source in promoting the diamond growth.

During the CVD process, D films grew with a columnar microstructure where the lateral grain size increased as the film became thicker as explain before. 289 To further understand the effects of UV laser wavelength on the lateral grain evolution, the cross-sectional microstructures of the D films were studied using SEM and are shown in Figure 6.4-6.

The UV laser in both cases induce a faster lateral grain size evolution which can be attributed to the suppression of secondary nucleation sites in the early stage of the D deposition. Optical images show an increase in the OH, CH, and C2 abundance with UV laser irradiations. All three species were reported to positively influence the diamond growth.

The OH radical functions as carbon etchant by etching the surface-bond hydrogen and stabilizing the sp 3 hybridized surface carbon bonds. 302 The CH radical was also believed to be helpful in D growth. 303 The addition of a C2 to a hydrogen-terminated diamond (110)

surface is energetically favorable and makes the growth of the existing crystal proceeds readily. 284,285 The increments of all three radicals explain why both growth rate and crystal quality were improved in the UV-laser-assisted combustion synthesis of diamond. To (b) C2H4. [304][305][306] Emissions from both C2 (swan band) and CH (A-X) were detected in their experiments. 308 These studies well support the observation of UV-laser-induced generation of a large quantity of C2 and CH in our OES study. 305,306,309,310 The 193 nm photodissociation of C2H4 has also been well characterized. Next, the effect of the excimer laser wavelength on the flame chemistry and D growth was studied. It was found that the laser wavelength was critical to the growth of high-quality D. A lower UV wavelength enhanced D growth and quality due to a higher photon energy, thereby promoting the photolysis of reactive species in the combustion flame.

Future work

In this work, the 3D laser printing of Cu/D composites was investigated, the first report on these materials. Several challenges were addressed to achieve this goal. Nevertheless, this research must move forward to provide more significant results. Also, each fabrication method has its own limitations on the reinforcement concentration.

It is crucial to know and understand the limitations of 3D laser printing in the manufacturing of Cu/D composite materials.

Another study which is vital to the development of 3D laser printing of Cu/D is to understand the role of the interphase (i.e., coating of D particles) on the wettability, consolidation, and final properties of the composites. For example, in Chapter 3, a multilayer coating was demonstrated, which had a Cr3C2 layer on the outside. This carbide is known to possess an enhanced wettability with molten Cu compared to Ti-based interphase. Thus, it will be interesting to study the effect of different coated D particles on the evolution of the printing parameters, densification, and final properties.

Also, in this work, the D is coated before printing, thus adding a step to the process. The in-situ formation of interphase could allow Cu/D composite to be 3D laser printed more efficiently and cost-effectively. For that, a low melting point, carbide-forming metal can be added to the powder mixture, such as Al. The short melting time provided by the SLM process may limit the reaction between Cu and Al while enabling the formation of an Al carbide interphase on the D particles.

Finally, the TC of the D film grown with and without a laser must be measured to confirm the advantage of laser-assisted CVD for producing a diamond film heat spreader. In summary, the work presented in this dissertation establishes a baseline for future development of 3D laser printing of highly conductive MMCs and heat spreaders for thermal management applications.

  de la fréquence de travail des dispositifs à base de semi-conducteurs avec leur miniaturisation a conduit à une surchauffe sévère, qui affecte leur durée de vie et leur fiabilité. Par conséquent, la gestion thermique est devenue une préoccupation importante dans le domaine microélectronique et doit être abordée. Le diamant (D) est connu pour être un excellent matériau pour la dissipation thermique car il possède l'une des conductivités thermiques les plus élevées de tous les matériaux naturels et possède une résistivité électrique élevée. D peut refroidir les puces électroniques de deux manières. Lorsqu'il est utilisé sous forme de film, D agit comme un diffuseur de chaleur. Lorsqu'ils sont utilisés sous forme de poudre, les Ds peuvent être introduits dans les métaux pour améliorer leur conductivités thermiques (TC) et apporter une stabilité dimensionnelle à des températures élevées. Les matériaux composites métal/D résultants sont ainsi d'excellents composants pour former des dissipateurs thermiques. Malgré l'attrait des matériaux à base de D en termes de performances thermiques, ils présentent souvent une géométrie simple, principalement en raison de la complexité d'usiner des matériaux à base de D dans des formes complexes. Dans cette étude, l'impression 3D laser de matériaux composites cuivre/D est proposée pour fabriquer des structures complexes de Cu/D qui pourraient remodeler leurs applications. Premièrement, la fabrication additive de Cu pur est optimisée et caractérisée. Puis, faute d'une affinité chimique entre Cu et D, une interphase est introduite dans le matériau composite. Plus tard, un procédé de revêtement de sel fondu est étudié pour produire un revêtement gradué et multicouche d'oxyde / carbure et de carbure / carbure, respectivement, sur des matériaux carbones. Ensuite, la fabrication additive de structures composites Cu/D est présentée. Enfin, le dépôt des films D est réalisé avec une flamme oxyacétylénique assistée laser. Les effets de l'introduction de lasers ultraviolets dans la flamme sont caractérisés en termes de réaction chimique, de qualité du film D first to my French and American advisors, Dr. Jean-Francois Silvain from the ICMCB and Prof. Yongfeng Lu from the College of Engineering of the University of Nebraska-Lincoln. They both have been incredibly supportive, patient, and resourceful since the beginning of the program. Thank you very much for your time and your immense contribution to this work, but also for all your support, trust, and help outside of the scientific scope. I want to thank the Direction General de l'Armement (DGA), for their financial support at the University of Bordeaux and Dr. Bruno Mortaigne. I also would like to thank the director of the ICMCB, Dr. Mario Maglione, for welcoming me to the institute. I want to thank Professors, Natale Ianno, Dennis Alexander and Jian Wang at the University of Nebraska Lincoln, as well as Dr. Veronique Gauthier from the University of Poitiers, Dr. Catherine Debiemme-Chouvy from La Sorbonne, Prof. Hansang Kwon from the University of Tohoku, and Prof. Jean-Marc Heintz for the University of Bordeaux for being part of my supervisory committee and taking time to judge my thesis.

  dans les années 1980, 40 W/cm 2 dans les années 2000, et atteint maintenant souvent 150 W/cm 2 . De plus, la température d'une puce de silicium (Si) par exemple, doit être maintenue en dessous de ~ 150 C pour éviter l'instabilité et la défaillance. Il apparaît donc que la capacité de refroidir un appareil est essentielle pour déterminer sa durée de vie, sa fiabilité et son efficacité. En effet, plus de la moitié des pannes électroniques sont liées à la surchauffe, tandis que 6%, 19% et 20% sont attribuables à la poussière, à l'humidité et aux vibrations. Par conséquent, la gestion thermique est devenue un des facteurs limitants du développement de la puissance électronique. Pour répondre à ces défis, nous proposons d'utiliser l'impression laser tridimensionnelle (3D) pour fabriquer des composite Cu/D ainsi que des films de diamant (D). Cette thèse se concentre sur la fabrication de dissipateurs de chaleur à base de D. Les sujets suivants sont traités : 1) la fabrication additive de matériaux Cu et Cu/D et 2) le dépôt de films de D assisté par laser ultraviolet (UV). Cette thèse résulte d'une collaboration internationale dans le cadre d'un double doctorat. Les deux partenaires sont le laboratoire de nano-ingénierie assistée par laser (LANE), qui fait partie du département de génie électrique et informatique de l'Université du Nebraska-Lincoln (UNL) aux États-Unis et le groupe de recherche sur les composites et la céramique de l'Institute de chimie et de la matière condensée de Bordeaux (ICMCB) à l'Université de Bordeaux (UBX) en France. La thèse comprend sept chapitres. Dans le chapitre 1, les motivations et le plan de la thèse sont présentés. Dans le chapitre 2, l'état de l'art lié à aux travaux de thèse sont passés en revue. Il a été montré que l'absence d'affinité chimique entre Cu et D se traduit par des liaisons interfaciales faibles et poreuses, conduisant à un mauvais transfert thermique. L'importance d'avoir une liaison interfaciale robuste est discutée. Il est montré qu'une interphase chimique joue un rôle crucial dans les propriétés thermiques globales des composites. Par la suite, les méthodes de fabrication et de production de composite métalliques sont discutées. Il est montré que seules des géométries simples sont possible avec les méthodes v conventionnelles et des étapes de post-traitement sont nécessaires pour fabriquer des structures de conception complexe.

viFigure 1 .

 1 Figure 1. (a) Illustration du processus SLM pour imprimer des dissipateurs de chaleur

Figure 2 .

 2 Figure 2. Illustration du mouillage de Cu sur les CF avec différentes interphases : (a)

Figure 3 .Figure 4 .

 34 Figure 3. (a) Schéma des trois étapes principales d'un processus de revêtement au sel

Figure 5 (

 5 Figure 5 (a) Micrographies SEM de la surface et des morphologies en coupe

Figure 1 . 1 - 1 3 Figure 1 . 2 - 1 5 Figure 2 . 2 - 1 10 Figure 2 . 2 - 2 12 Figure 2 . 3 - 1 14 Figure 2 . 3 - 2 16 Figure 2 . 4 - 1 18 Figure 2 . 4 - 2 21 Figure 2 . 4 - 3 21 Figure 2 . 4 - 4 24 Figure 2 . 4 - 5 25 Figure 2 . 4 - 6 26 Figure 3 . 1 - 1 ( 31 Figure 3 . 2 - 1 ( 33 Figure 3 . 2 - 2 ( 35 Figure 3 . 3 - 1 37 Figure 3 . 3 - 3 40 Figure 3 . 4 - 1 ( 42 Figure 3 . 4 - 2 43 Figure 3 . 4 - 3 ( 45 Figure 3 . 4 - 5 ( 47 Figure 3 . 4 - 6 ( 49 Figure 3 . 4 - 7 50 Figure 3 . 4 - 8 ( 52 Figure 3 . 4 - 9 54 Figure 3 . 4 - 10 57 Figure 3 . 4 - 12 58 Figure 3 . 4 - 59 Figure 4 . 1 - 1 63 Figure 4 . 2 - 1 65 Figure 4 . 3 - 1 69 Figure 4 . 3 - 2 70 Figure 4 . 3 - 3 71 Figure 4 . 3 - 4 73 Figure 4 . 3 - 5 74 Figure 4 . 3 - 6

 111312152211022212231142321624118242212432124424245252462631131321333223533137333403414234243343453454734649347503485234954341057341258345941163421654316943270433714347343574436 Figure 1.1-1 Illustration of a multichip power electronics module. ....................................3 Figure 1.2-1 Schedule for the four-year dual Ph.D. program. ............................................5 Figure 2.2-1 Illustration of (a) the combustion flame setup and (b) IR laser vibrational excitation of C2H4 molecules. 14 .........................................................................................10 Figure 2.2-2 Scanning electron microscopy (SEM) micrograph of (a) D powder and (b) CVD diamond film. 14 .........................................................................................................12 Figure 2.3-1 SEM micrographs of the Cu-D interface. .....................................................14 Figure 2.3-2 Illustration of the powder metallurgy route (e.g., uniaxial hot pressing). ....16 Figure 2.4-1 Illustration of the SLM/SLS processes. ........................................................18 Figure 2.4-2 Illustration of LDED. ...................................................................................21 Figure 2.4-3 Illustration of the different wetting behaviors. .............................................21 Figure 2.4-4 Photograph of aluminum silicon (Al-Si) 10 mg lattice structures, reported by K.K Wong et al. 63 ...............................................................................................................24 Figure 2.4-5 Photographs of 3D-printed heat sinks made of Al-Si10Mg, by M Wong et al. .66 ....................................................................................................................................25 Figure 2.4-6 Photograph of Cu coil printed by C. Silbernagel et al. 74 ..............................26 Figure 3.1-1 (a) Schematic illustration of the SLM process to print Cu heat sinks, (b) photo of a printed heat sink on an electronic chip with the thermal image in the insert. ............31 Figure 3.2-1 (a) Schematic of the homemade SLS system, (b) line scanning strategy of the laser sintering, and (c) temperature profile of the sintering pool. ......................................33 Figure 3.2-2 (a) SLM parameters used in this study, (b) SEM micrographs of the copper powders, and (c) schematic illustration of the stripe scan strategy....................................35 Figure 3.3-1 Powder metallurgy steps to obtain dense Cu parts: (a) raw powders, (b) after cold pressing, and (c) after the SLS process. .....................................................................36 Figure 3.3-2 (a) SEM micrographs and (b) 3D surface mapping of Cu layer fabricated using SLS from loose and cold-pressed Cu powders, respectively (c) XRD patterns of the raw and sintered powders...................................................................................................37 Figure 3.3-3 SEM micrographs of Cu sample surfaces fabricated by SLS at different scan speeds: (a) surface morphologies, (b) Cu part cross-sectional overviews, and (c) crosssectional microstructures. ..................................................................................................40 Figure 3.4-1 (a) Photo of printed Cu cubes with laser scan speeds from 200 -600 m/s and hatch distances from 0.08 to 0.12 mm, respectively. (b) 2D map of the printing quality with respect to the scan speed and the hatch distance................................................................42 Figure 3.4-2 XRD diffractograms of the printed part (top) and Cu powders (bottom). ..43 Figure 3.4-3 (a) and (b) SEM micrographs of Cu powder and part, (c) and (d) EDX point spectra and (e) summarized EDX analyses. .......................................................................44 Figure 3.4-4 (a) 2D map of the part density, and (b) the average surface roughness as a function of the scan speed and the hatch distance. ............................................................45

Figure 4 . 3 - 7 76 Figure 4 . 3 - 8 79 Figure 4 . 3 - 9 80 Figure 4 . 3 - 81 Figure 4 . 3 - 11 86 Figure 4 . 3 - 88 Figure 4 . 3 - 90 Figure 4 . 3 - 16 91 Figure 4 . 3 - 17 A 91 Figure 4 . 3 - 18 92 Figure 4 . 4 - 1 95 Figure 4 . 4 - 2 97 xixFigure 4 . 4 - 3 99 Figure 4 . 4 - 4 Figure 4 . 4 - 7 Figure 4 . 4 - 8 Figure 5 . 2 - 3 Figure 5 . 3 - 1 Figure 5 . 3 - 3 Figure 5 . 4 - 2 Figure 5 . 4 - 3 Figure 5 . 4 - 4 Figure 5 . 4 - 5 Figure 5 . 4 - 6 Figure 6 . 1 - 1 Figure 6 . 2 - 1 Figure 6 . 2 - 2 Figure 6 . 3 - 2 Figure 6 . 3 - 3 Figure 6 . 3 - 4 Figure 6 . 3 - 5 (Figure 6 . 3 - 6 Figure 6 . 3 - 7 Figure 6 . 3 -Figure 6 . 4 - 1 Figure 6 . 4 - 2 Figure 6 . 4 - 3 Figure 6 . 4 - 4 Figure 6 . 4 - 6 Figure 6 . 4 - 8 Figure 6 . 4 - 9 Figure 6 . 4 -Figure 6 . 4 - 13

 43776438794398043814311864388439043169143179143189244195442974439944444744852353153354254354454554661162162263263363463563663763641642643644646648649646413 Figure 4.3-7 XPS survey and high-resolution spectra of the C 1s, Ti 2p, and O 1s peaks for (a) TiOy-TiCx and (b) TiCx CFs, respectively............................................................76 Figure 4.3-8 SEM micrographs of uncoated CFs after Cu wetting at 1200 ºC for 10 min. ............................................................................................................................................79 Figure 4.3-9 Cu contact angles measured on CFs coated with (a) TiCx and (b) TiOy-TiCx prepared at 800 and 950 ℃, respectively. (c) Adhesive (wetting) energy vs. contact angle measured in this work (•) compared to values reported by others for Cu on CF (■), 49 TiC (♦), 50 and TiO (▲). 51 ..........................................................................................................80 Figure 4.3-10 SEM micrographs showing the wetting behavior of Cu on the coated CFs prepared at (a), (b) 800, and (c), (d) 950 ℃. (e) Measured Cu area coverage with the equivalent TiOy content and droplet diameter as functions of the contact angle. .............81 Figure 4.3-11 Schematic of the 3 main steps of the molten-salt coating process. ............82 Figure 4.3-12 (a) SEM micrographs of rapidly quenched salt at 800 after 0.5 h, XRD diffractograms of washed CFs for coating times ranging from 0.5 to 5 h for (b) 800 and (c) 950 ºC, (d) XRD peak intensities as functions of the coating time. ..................................85 Figure 4.3-13 SEM micrographs of the cross-sectional views of the CFs coated at different temperatures and times used to determine the diffusion coefficient and activation energy. ............................................................................................................................................86 Figure 4.3-14 (a) Plots of the coating thickness as functions of the square root of time for different coating temperatures; (b) Arrhenius plot of Deff as a function of the coating temperature; (c) a schematic of the counter-diffusion; cross-sectional view of coated CFs prepared at (d) 800 and (e) 950 ℃, (f) EBSD mapping of the cross-section of a TiC-covered CF showing the distribution of grain orientations. ............................................................88 Figure 4.3-15 SEM micrographs of (a) an as-received D particle, (b) a D particle with TiO2/TiC coating; high magnification SEM micrographs of D surface coated at (c) 800, (d) 850, (e) 900, and (f) 950 ºC, respectively. .........................................................................90 Figure 4.3-16 Surface analysis of diamond coated with TiO2-TiC at (a) 800, (b) 850, (c) 900, and (d) 950 ºC; (e) the curve for the average surface roughness vs. the salt-bath temperature. .......................................................................................................................91 Figure 4.3-17 A typical Raman spectrum of TiO2-TiC coating. .......................................91 Figure 4.3-18 Raman mapping of diamond coated with TiO2-TiC at (a) 800, (b) 850, (c) 900, and (d) 950 ºC, respectively. The red and green color represent TiO2 and TiC, respectively. .......................................................................................................................92 Figure 4.4-1 SEM micrographs of coatings on CFs with M = Ti, Cr, and Ti + Cr. (a) crosssectional views, (b) schematic illustrations of the coatings with M = Ti, Cr, and Cr + Ti, and (c) XRD diffraction patterns of TiC, Cr3C2, and TiC + Cr3C2 coatings, and (d) AES depth profiles of the CFs coated in a Ti + Cr salt solution. ...............................................95 Figure 4.4-2 The steps to form the multilayer coating: (a)-(d) SEM micrographs after salt quenching at different times of 0, 10, 20, and 30 min, (e) EDX analysis of the salt quenched at 10 min, and (f) schematic illustration of the coating steps. ...........................................97

Figure 1 . 1 - 1

 111 Figure 1.1-1 Illustration of a multichip power electronics module.

  program. The two partners are the Laser-Assisted Nano-Engineering (LANE) lab, part of the Department of Electrical and Computer Engineering at the University of Nebraska-Lincoln (UNL) in the U.S. and the Composite and Ceramics research group at the Institute of Condensed Matter Chemistry of Bordeaux (ICMCB) at the University of Bordeaux (UBX) in France. My work in the U.S. and France concerned 3D laser printing of Cu/D composites, processing of graded and multilayer coatings on carbon-based materials for Cu/D composites, and deposition of diamond films. The supervisors for these research activities were Professor Yongfeng Lu at UNL and Professor Jean-François Silvain at ICMCB/UBX. This four-year program was divided into two years in the U.S. and two years in France, as shown in Figure 1.2-1.

Figure 1 . 2 - 1

 121 Figure 1.2-1 Schedule for the four-year dual Ph.D. program.

Figure 2 . 2 - 1

 221 Figure 2.2-1 Illustration of (a) the combustion flame setup and (b) IR laser vibrational excitation of C2H4 molecules. 10

Figure 2 . 2 - 2

 222 Figure 2.2-2 Scanning electron microscopy (SEM) micrograph of (a) D powder and (b) CVD diamond film. 10

2 . 3 - 1 .

 231 Several studies report low TCs (> 230 W/m.K) for Cu/D composites. The degradation of the TC is due to the weak affinity between Cu and D and the creation of pores at the interface, as shown in Figure Thus, the synthesis of the interphase between the Cu and C reinforcement is essential to achieve high thermal conductivity.[START_REF] Schubert | Interfacial design of Cubased composites prepared by powder metallurgy for heat sink applications[END_REF][START_REF] Tao | Properties and microstructure of Cu/diamond composites prepared by spark plasma sintering method[END_REF]

Figure 2 . 3 - 1

 231 Figure 2.3-1 SEM micrographs of the Cu-D interface.

2. 3 . 2

 32 In situ formation of the interphase between Cu and D Adding an alloying element was considered to improve the wettability of Cu and C during the sintering of the composite materials. Mortimer et al. studied the wettability of Cu and Cu alloys on C and showed improvement in the contact angle compared to pure

  et al. studied the formation of molybdenum carbide (Mo2C) and titanium carbide (TiC) via a molten-salt process on carbon fibers (CFs). 46 They demonstrated the fabrication of a homogenously coated layer. Next, the manufacturing of Cu/CF composite was performed via spark plasma sintering (SPS). The TC of the Cu/CF without interphase was measured to be 206 W/m.K. However, with Mo2C and TiC, the TC increased up to 290 and 297 W/m.K, respectively. Li et al. manufactured Cu/D composites via a gas pressure infiltration with a Ti-based coating on D particles. A TC of 716 W/m.K was achieved with a D volume fraction of 65 vol.% and a size of about 70 μm. 47 Kang et al. demonstrated the deposition of Cr3C2 on D particles via a molten-salt process. The resulting Cu/D prepared by vacuum pressure infiltration containing 65 vol.% of D (70 μm) exhibited a thermal conductivity of 562 W/m.K.

Figure 2 . 3 - 2 .

 232 The pressure enhances the powder packing density and creates dislocations in the metal powders, helping the densification. In these methods, the powders are placed into a mold (usually graphite). Thus, the mold design defines the final shape of the MMCs. The geometry is usually simple (i.e., cubes, cylinder) to obtain a homogenous pressure during the sintering. Subtractive manufacturing is widely employed to fabricate more sophisticated shapes.[START_REF] Guillemet | Diamond-based heat spreaders for power electronic packaging applications[END_REF]

Figure 2 . 3 - 2

 232 Figure 2.3-2 Illustration of the powder metallurgy route (e.g., uniaxial hot pressing).

Figure 2 . 4 - 1

 241 Figure 2.4-1 Illustration of the SLM/SLS processes.

Figure 2 . 4 - 2

 242 Figure 2.4-2 Illustration of LDED.

Figure 2 . 4 - 3

 243 Figure 2.4-3 Illustration of the different wetting behaviors.

69

 69

Figure 2 . 4 - 4

 244 Figure 2.4-4 Photograph of aluminum silicon (Al-Si) 10 mg lattice structures, reported by K.K Wong et al.[START_REF] Wong | Saturated pool boiling enhancement using porous lattice structures produced by Selective Laser Melting[END_REF]

Figure 2 . 4 - 5

 245 Figure 2.4-5 Photographs of 3D-printed heat sinks made of Al-Si10Mg, by M Wong et al..66

  manufacturing of 3D objects. C.Silbernagel et al. studied the SLM of intricate Cu parts using, as shown in Figure2.4-6, a moderated laser power. In this study, a laser power of 200 W was used to print complex Cu coils. However, the low laser power resulted in the fabrication of porous Cu parts (density = 85%). Thus the electrical conductivity (i.e., TC) of the printed part was severely degraded and only reached about 50% of the International Annealed Copper Standard (IACS).[START_REF] Silbernagel | Electrical resistivity of pure copper processed by mediumpowered laser powder bed fusion additive manufacturing for use in electromagnetic applications[END_REF]

Figure 2 . 4 - 6

 246 Figure 2.4-6 Photograph of Cu coil printed by C. Silbernagel et al.88

Figure 3 . 1 - 1 Figure 3

 3113 Figure 3.1-1 (a) Schematic illustration of the SLM process to print Cu heat sinks, (b)

Figure 3

 3 Figure 3.2-1 (a) schematically shows the home-built SLS setup used in this work. A

  printing parameters are summarized in Figure 3.2-2 (a). The printer is equipped with a CW fiber laser (IPG photonics, YLM, Yb:YAG, λ = 1070 nm). The maximum output power is 400 W with a spot diameter of 70 μm on the powder-bed surface. Spherical Cu powders with an average size of 45 μm were used in this study (US Nano, Cu 99.5%). The size of the powder and dispersity of the powder were determine based on the powder flowability (spherical), bed compaction (high dispersity), and requirements of the SLM 125 printer (10 to 50 μm). A SEM micrograph of the powders is shown in Figure 3.2-2 (b).The powder was stored under agon to prevent surface oxidation. The printing chamber was filled with argon (Ar)(99.9%, Matheson) to reduce an oxygen level below 0.09%. A constant pressure of 18 mbar was maintained during all experiments. The printing parameters were optimized by adjusting the hatch distance from 0.08 to 0.12 mm, and the scan speed from 200 to 600 mm/s. The laser power and layer thickness were kept constant at 400 W and 0.03 mm, respectively, throughout the study. A stripe scan strategy was used in this study, as illustrated in Figure3.2-2 (c). The structures printed were designed using SolidWork ® . Magics ® software was used to prepare print jobs and establish the printing parameters. The parts were printed on a stainless-steel building plate (stainless steel, 123

Figure 3 .

 3 Figure 3.2-2 (a) SLM parameters used in this study, (b) SEM micrographs of the copper

Figure 3 . 3 - 1 .

 331 Both a loosely packedCu powder bed and a cold-pressed Cu pellet were sintered using an Nd:YAG laser at a constant laser scan speed of 125 µm/s and laser power of 5 W. The density of the loose powder and the cold-pressed pellet were estimated to be about 30 % and 63 %, respectively.

Figure 3 . 3 - 1

 331 Figure 3.3-1 Powder metallurgy steps to obtain dense Cu parts: (a) raw powders, (b)

Figure 3 .

 3 Figure 3.3-2 (a) shows the SEM micrographs of parts fabricated from the loose-powder

Figure 3 .

 3 3-1 (a) and (b).

Figure 3 .

 3 Figure 3.3-2 (a) SEM micrographs and (b) 3D surface mapping of Cu layer fabricated

Figure 3 .

 3 Figure 3.3-2 (b), sintering of the loose-powder bed resulted in a rough surface containing

Figure 3 . 3 - 3

 333 shows the surface morphologies and the cross-sectional micrographs of 50-μm-thick parts fabricated by SLS from cold-pressed Cu pellets at different scan speeds. As shown in Figure 3.3-3 (a), smooth surfaces without balling globules were obtained for all cases and the number of pores decreased as the laser scan speed decreased. It was observed in Figure 3.3-3 (b) that homogenous sintering of the Cu powders was obtained throughout the whole thickness for each scan speed. Moreover, the cross-sectional micrographs, shown in Figure 3.3-3 (c), revealed the consolidation behaviors at different scan speeds. At a scan speed of 125 µm/s, necks formed between Cu particles. By reducing the scan speed to 62 µm/s, the sintering of the powders was enhanced. At a scan speed of 25 µm/s, complete densification of the Cu powders was attained. The surface of the parts shows dimple-like features, which are a typical characteristic of ductile fracture.106,119 As the scan speed decreased, the sintering time increased, which means more energy was delivered and absorbed by the powders and thus the atomic diffusion was promoted, which led to denser parts.120

Figure 3 . 3 - 3

 333 Figure 3.3-3 SEM micrographs of Cu sample surfaces fabricated by SLS at different scan

Figure 3

 3 show a photo of the printed parts on the base plate and a 2D map depicting the quality of the parts at different scan speeds and hatch distances-cracked, high quality, and porous samples are represented by green squares, red circles, and blue triangles, respectively. A low scan speed (200 to 350 mm/s) resulted in cracked samples for all hatch distances studied. With a low scan speed, the energy deposited greatly exceeded the amount required to melt the Cu powders. The high temperature of the molten pool generated a large thermal gradient in the solid Cu, which ultimately caused interlayer delamination and debonding from the build plate.[121][122][123] Note that, at a scan speed of 200 mm/s (Figure3.4-1 (a)), the print was stopped because bending/debonding of Cu from the build plate caused severe damage to the powder recoater. Printing with a higher scan speed of 400 mm/s and a hatch distance of 0.12 mm produced smooth Cu parts without defects. Similar results were also obtained with scan speeds ranging from 450 to 500 mm/s and hatch distances from 0.08 to 0.09 mm (Figure3.4-1 (a)-(b). Finally, scan speeds exceeding 500 mm/s yield rough and porous samples for the majority of the hatch distances investigated (Figure 3.4-1 (a)-(b)). When insufficient laser energy is applied partial melting and sintering of the powders occurs and subsequently resulting in pore generation. 124

Figure 3 .

 3 Figure 3.4-1 (a) Photo of printed Cu cubes with laser scan speeds from 200 -600 m/s

Figure 3 . 4 - 2

 342 Figure 3.4-2 XRD diffractograms of the printed part (top) and Cu powders (bottom).

Figure 3 .

 3 4-3 (a) and (b) present the SEM micrographs of the Cu powders and polish Cu parts (printing parameters: power = 400 W, layer thickness = 0.03 mm, scan speed = 400 mm/s, and hatch distance = 0.12 mm), respectively.Figure 3.4-3 (c) and (d) show the EDX point spectra of the Cu powders and parts, respectively.

Figure 3 .

 3 Figure 3.4-3 (e) summarizes the EDX analysis where it can be observed that the powders

Figure 3 .

 3 Figure 3.4-3 (a) and (b) SEM micrographs of Cu powder and part, (c) and (d) EDX point

Figure 3 .

 3 Figure 3.4-4 (a) 2D map of the part density, and (b) the average surface roughness as a

  A slow laser scanning speed generates a significant temperature gradient in the molten pool and induces a great difference in the surface tension between the edge and the center of the molten pool. As a result, the molten Cu tends to flow from the hot molten pool towards the colder edges of the pool, thus creating protrusion at the pool borders (Figure 3.4-5 (c)). 130 At 400 mm/s, the part shows a smooth and flat surface finish with minimal defects (Figure 3.4-5 (d) and (e)). At a moderate scan speed, the temperature of molten Cu is sufficiently high to achieve a desirable metal viscosity, resulting in a flat and stable molten pool, as illustrated in Figure 3.4-5 (f). At 600 mm/s, the surface appears to be porous with the balling effect (Figure 3.4-5 (g) and (h)). The insufficient laser energy deposition that occurs at a fast scan speed leads to an unstable molten pool (i.e., fluctuation), with poorly melted powders and a slow molten metal flow, leading to the formation of large pores (> 80 μm) (Figure 3.4-5 (i) and (f)) and bumps on the surface (i.e., the balling effect). 131

Figure 3 . 4 - 5 (

 345 Figure 3.4-5 (a), (d), and (g) SEM micrographs of as-printed surfaces and (b), (e), and

Figure 3 .

 3 4-6 (a).Figure 3.4-6 (b) lists the parameters of the laser beam and are conformed to the EN ISO 11146-2 measurementmethod. The beam quality is evaluated by the beam quality factor (BPP) and the effective beam quality factor (M 2 ). It is measured that the BPP has a value of 0.396 mm/rad and an M 2 of 1.16. The results imply an excellent beam quality when focused to a small spot with the conservation of the original output power.138,139

Figure 3 . 4 - 6 (

 346 Figure 3.4-6 (a) Laser beam profile, and (b) beam characteristics, measured inside the

3 . 4 - 7 .

 347 structures has the potential to significantly enhance the functionality of current heat sinks to cool electronic devices. The full potential of the SLM method was investigated by fabricating different heat sink structures, as presented in Figure The first heat sink is a commercial (COM) columnar type.

Figure 3 . 4 - 7

 347 Figure 3.4-7 Top-and side-view optical and SEM micrographs of (a) a commercial

Figure 3 . 4 - 7 (

 347 Figure 3.4-7 (a) and (b) show optical and SEM micrographs of COM and printed heat

Figure 3

 3 of the sample preparation for the four-point probe and hardness measurements. The heat sink pillars were cut from the base plate for electrical conductivity measurements, while the base plate was used for hardness tests.

Figure 3 . 4 - 8 (

 348 Figure 3.4-8 (a) Schematic illustration of the sample preparation for measurements of the

( 2 .

 2 2 x 10 -8 WΩ/K 2 ). The TC of the COM and the printed heat sinks was 378 and W/m.K, respectively. The TC of the commercial and printed heat sink measured in this study are comparable to the one measured by C. Vincent et al. on pure Cu prepared by a powder metallurgic route with similar density. 142 Note that the incertitude in the electrical conductivity measurements is derivate from the incertitude on the pillar dimensions measurement and the resistivity analysis.

Figure 3 . 4 - 9 (

 349 Figure 3.4-9 (a) shows the SEM micrographs of the indentation of the COM and printed

Figure 3 . 4 - 9

 349 Figure 3.4-9 SEM micrographs of the base plates of the commercial (COM) and printed

Figure 3 . 4 -

 34 10 shows the experimental set-up where the heat sink was attached to a memory card chip placed on a hot plate heated to 100, 150, and 200 ºC, respectively. A fan was located 10 cm from the top of the heat sink to produce an airflow. The temperature of the electronic chip was measured using an IR camera about 0.5 m from the heat sink.

Figure 3 . 4 - 10 3

 34103 Figure 3.4-10 Schematic illustration of the experimental set-up to measure the heat

Figure 3 . 4 -

 34 Figure 3.4-11 IR images of chip alone, and with COM, printed, helix, and tubes heat

Figure 3 . 4 - 12

 3412 Figure 3.4-12 Chip temperature vs. time plots for different heat sink structures heated at

Figure 3 . 4 -

 34 Figure 3.4-13 (a) Cooling effect (Cooling = Tchip -Tchip with heat sink) of different heat sinks

Figure 4 . 1 - 1

 411 Figure 4.1-1 Schematic of a complex multilayer coating formed in a single-step, molten

Figure 4 . 2 - 1 .

 421 Figure 4.2-1. First, CFs (Mitsubishi Chemical DIALEAD ™ K223HM) or D particles

Figure 4 . 2 - 1

 421 Figure 4.2-1 Illustration of the molten-salt process for preparing TiC-based coatings

Figure 4 . 3 - 1 .

 431 In this study, the reaction time was 5 h at temperatures ranging from 800 to 950 ºC. The SEM micrographs of the coatings display a drastic change in the morphology with different temperatures. Specifically, at 800 and 850 ºC, the surfaces appeared to be a porous collection of particle-like features with open, interconnected porosity of more than 50% (pore sizes ranging from 300 to 1000 nm). This suggests that the growth was from reactive sites on the C surface with little consolidation by surfacediffusion-driven sintering. Both Xiaoguang Liu et al.184 and Qian Liu et al.186 observed similar porous coatings formed by Ti reaction with C substrates at 750 ºC in a potassium chloride : lithium chloride (KCl:LiCl) salt and at 900 ºC in a lithium chloride : potassium chloride : potassium fluoride (LiCl:KCl:KF) salt, respectively.

Figure 4 . 3 - 1

 431 Figure 4.3-1 SEM micrographs of the Influence of the coating temperature, ranging from

Figure 4 . 3 - 2 ).

 432 It can be seen that for each temperature and a coating time of 5 h, diffraction peaks at 2 =35.9°, 41.7°, 60.5°, 72.4°, and 76.1º corresponded to the (111), (200), (220), (311), and (222) planes of a facecentered cubic (FCC) titanium carbide (TiC) phase. The strong and sharp (002) diffraction peak at 2 = 26.4° is assimilated to CFs and confirmed its high degree of graphitic crystallinity.

Figure 4 . 3 - 2 4 . 3 - 3 .

 432433 Figure 4.3-2 XRD diffractogram of the CFs coated with TiCx at different temperatures

Figure 4 . 3 - 3

 433 Figure 4.3-3 Influence of the coating temperature, ranging from 800 to 950 ºC, on the

Figure 4 . 3 - 1 4 . 3 - 1 ) 3 - 4 )

 43143134 was Ti:C = 0.125. The fully dense (0% porosity) coating thickness was predicted to be 0.55 μm via the mass balance calculation, suggesting an outer fiber diameter of 11.45 μm. This is in reasonable agreement with the outer diameter of 11.6 μm observed for CFs reacted at 950 ºC (Figure and supported the observations that the porosity was low (<10%). Similarly, the coating prepared at 800 ºC has a measured outer diameter of ~13.5 μm, suggesting a porosity of ~65%, again in agreement with the observations. The relationships between the TiCx lattice parameters and the coating temperature were determined from the XRD data using the Le-Bail method. 185 Figure 4.3-4 represents the refinement results where the black, red, and blue lines correspond to the experimental, refined diffractogram, and residual intensities, respectively. The Bragg factor expresses the agreement between the calculated intensities and the measured ones (insert in Figure 4., which confirms an excellent agreement between the experimental and refined diffractograms.

Figure 4 . 3 - 4

 434 Figure 4.3-4 Le-Bail-structural refinement using FullProf Suite for XRD diffractograms

Figure 4 . 3 - 5

 435 Figure 4.3-5 Influence of the coating temperature, ranging from 800 to 950 ºC, on the

Figure 4 . 3 - 6

 436 shows the presence of not only TiCx on the CF surfaces but also an outer layer of TiOy. The elemental analyses at different depths of the coatings prepared at 800 ºC show a thin TiOy outer layer that extends to ~20% of the coating thickness. Based on the results in Figure 4.3-6 (a) and (c), the thickness of the coating at 800 ºC is about 1.2 µm, suggesting an partially oxidized outer layer of 250 nm. In contrast, the coatings prepared at 950 ºC were significantly denser with a thickness of ~0.30 µm (Figure 4.3-6 (b) and (d)) and a TiOy outer layer of < 5 nm. In other words, the 950 ºC coating is essentially 100% TiCx.

Figure 4 . 3 - 6

 436 Figure 4.3-6 Surface morphology, cross-sectional and AES maps of the coated CFs

Figure 4 . 3 - 7

 437 Figure 4.3-7 XPS survey and high-resolution spectra of the C 1s, Ti 2p, and O 1s peaks

Figure 4

 4 2 (𝑇 -1085) for Cu, 194 and  is the equilibrium contact angle. The results show that Cu does not wet CFs (

Figure 4 . 3 - 8 Figure 4

 4384 Figure 4.3-8 SEM micrographs of uncoated CFs after Cu wetting at 1200 ºC for 10

Figure 4 . 3 - 9

 439 Figure 4.3-9 Cu contact angles measured on CFs coated with (a) TiCx and (b) TiOy-

Figure 4 . 3 - 4 . 3 . 4

 43434 Figure 4.3-10 SEM micrographs showing the wetting behavior of Cu on the coated CFs

Figure 4 . 3 - 11

 4311 Figure 4.3-11 Schematic of the 3 main steps of the molten-salt coating process.

4 . 3 -

 43 transport of C. Semi-qualitative analyses of XRD peaks intensities of coated CFs were used to follow the reaction versus time at temperatures of 800 and 950 ºC as presented in Figure 12 (b) and (c), respectively. The integrated XRD peak intensities at 800 ºC show significant formation of TiCx even after 0.5 h (Figure 4.3-12 (d)). Moreover, the amount of TiC continuously increases with time as Ti decreases. After 3 h of reaction, the Ti isonly present in trace amounts, and none is observed after 5 h. At 950 ºC, complete Ti dissolution occurs essentially only after 0.5 h as evidenced by the lack of a Ti XRD peak and the high TiCx signal. Also, the TiC peak intensity is essentially constant from 0.5 h to 5h within the XRD analysis error limits. These results agree with our calculations for the rapid dissolution and diffusion of Ti particles in molten KCl.

Figure 4 . 3 -

 43 Figure 4.3-12 (a) SEM micrographs of rapidly quenched salt at 800 after 0.5 h, XRD

Figure 4 . 3 -

 43 Figure 4.3-13 SEM micrographs of the cross-sectional views of the CFs coated at

Figure 4 . 3 -

 43 14 (a), the coating thickness (x) shows a linear increase with the square root of the coating time. The diffusivity, Deff, can be retrieved from the slope of the curve. The value increases with the coating temperature ranging from 3.8×10 -13 to 1.3×10 -12 cm 2 /s. These values, agree reasonably with values report by Koichiro Koyama et al. for the diffusion of C in TiC containing a small amount of oxygen and over a temperature range similar to that used in this work.203

Figure 4 . 3 -

 43 14 (b).As might be expected, the activation energy is lower than the one reported by Knuth Albertsen et al.for grain boundary diffusion of C in bulk TiC below 1100 ºC (174 kJ/mol).206 This could be due to the enhancement of the grain boundary diffusion due to the difference in microstructures between bulk and coated TiC.206,207

Figure 4 . 3 -

 43 Figure 4.3-14 (a) Plots of the coating thickness as functions of the square root of time for

Figure 4 . 3 -

 43 Figure 4.3-15 (a) shows that the as-received D powder has an average diameter of 50 µm

Figure 4 . 3 -

 43 Figure 4.3-15 SEM micrographs of (a) an as-received D particle, (b) a D particle with

Figure 4 . 3 - 16

 4316 Figure 4.3-16 Surface analysis of diamond coated with TiO2-TiC at (a) 800, (b) 850, (c)

Figure 4 . 3 - 17 AFigure 4 . 3 - 18 )

 43174318 Figure 4.3-17 A typical Raman spectrum of TiO2-TiC coating.

Figure 4 . 3 - 18

 4318 Figure 4.3-18 Raman mapping of diamond coated with TiO2-TiC at (a) 800, (b) 850,

  result of oxidation of reactive Ti sites near the exterior coated surfaces. The oxide-rich carbide coating shows significantly improved wettability with molten Cu compared to the carbide-rich layers. Notably, the Cu wetting angle for the TiOy-TiCx-CF samples prepared at 800 ºC is ~80º ± 5º with a Cu surface coverage of ~50% versus ~115º and ~ 10% for those of the TiCx-CF samples prepare at 950 ºC. CFs alone show mostly no Cu coverage.By varying the reaction temperature (800 to 950 ºC) and time (1 to 5 h) and using rapid thermal quenching, the formation mechanism is investigated. The results indicate a threestep process: 1) a short induction period controlled by rapid Ti dissolution in the molten salt, 2) rapid Ti mass transport in the molten salt and fast reaction at the CF surfaces, and 3) mass transport through the coating layer that is controlled by a competition of Ti and C diffusions.Finally, the graded TiO2-TiC layers developed on CFs are applied on D particles. It is shown that similar results are obtained. A TiO2 layer forms on the D surfaces. The concentration of Ti oxide is found to be related to the coating temperature. A lower coating temperature leads to a lower diffusion of the reactant resulting in a partial oxidize coating.

Figure 4 . 4 - 1 (

 441 a) shows the CFs coated using M = Ti, M = Cr, and M = Ti + Cr, revealing homogenous layers on all CFs throughout the anisotropic structure. The thicknesses of the 3 layers were similar, about 700 nm (Figure 4.4-1 (a)). However, the coating obtained by a Ti and Cr powder mixture consisted of two layers of contrast, suggesting the formation of two phases, as illustrated in Figure 4.4-1 (b).

Figure 4 . 4 - 1

 441 Figure 4.4-1 SEM micrographs of coatings on CFs with M = Ti, Cr, and Ti + Cr. (a)

Figure 4 . 4 - 2 (

 442 Figure 4.4-2 (a) shows the SEM micrographs of the salt quenched in the early stage of

Figure 4 . 4 - 2 4 . 4 - 3 )

 442443 Figure 4.4-2 The steps to form the multilayer coating: (a)-(d) SEM micrographs after salt

Figure 4 . 4 - 2 (

 442 c) shows the salt mixture quenched after 20 min. At this point, the Ti particles and the shells around the Cr particles were no longer observable, suggesting their complete reaction. After 30 min, Cr powders dissolved in the salt, forming the outer Cr3C2 layer (Figure 4.4-2(d)). The previous observations show an oscillatory reaction behavior in the molten salt, leading to the formation of a complex multilayer coating. Especially, in the early stages of the coating process, Ti and Cr saturated the liquid salt and the first Cr3C2 layer forms. Next, the TiC layer nucleated on the CFs and Ti surrounds the Cr particles, preventing further Cr dissolution. After the completion of the TiC layer, Cr dissolved in the salt reacted with the CFs, creating the final Cr3C2 layer, as illustrated in Figure 4.4-2 (f).

Figure 4 . 4 - 3

 443 Figure 4.4-3 XRD diffractogram of the quenched CFs after salt dissolution for different

Figure 4

 4 Figure 4.3-13 and Figure 4.4-4.

Figure 4 . 4 - 4

 444 Figure 4.4-4 BSE-mode SEM cross-sectional views of CFs coated with Cr3C2 for 1 to 5 h

Figure 4 .

 4 Figure 4.4-5 (a) C diffusion coefficients in both Ti and Cr as functions of the coating

Figure 4 . 4 - 6 (

 446 b) depicts the formation of pores at the TiC-CF interphase with a remaining CF diameter of 7 µm (~ 11 µm initially). In contrast, channels formed near the interface between the Cr3C2 layer and the CF surfaces, with its diameter reduced to 5 µm (Figure4.4-6 (c)). These findings imply that C has faster reaction and diffusion in Cr3C2 produced by molten salt than in TiC, which supports the kinetic analyses carried out in this study.

Figure 4 . 4 - 6 ( 4 . 4 . 4

 446444 Figure 4.4-6 (a) schematic illustration of the Kirkendall effect and coatings at 950 ºC for

Figure 4 .

 4 Figure 4.4-7 (a). Prior to the oxidation test, the surface of the TiC layer (M = Ti) appeared

Figure 4 . 4 - 7

 447 Figure 4.4-7 SEM micrographs before and after testing for CFs with (a) TiC, (b) Cr3C2,

Figure 4 . 4 - 8

 448 Figure 4.4-8 Raman spectra of original CFs and multilayer-coated CFs after the flame

Figure 5 . 1 - 1 (

 511 Figure 5.1-1 (a) schematic illustration of the laser directed energy deposition system

  A 5 vol.% of coated D was mixed with spherical Cu powder with a mean diameter of 93 μm (99.5% US Research Nanomaterials, Inc.) as shown in Figure 5.1-1 (b). AM was performed using an Optomec, Inc., LENS ® 3D Hybrid Machine Tool System (Optomec, Inc. LENS MR-7). The laser power (λ = 1064 nm, gaussian profile, continuous wave, max. power : 2 KW, spot diameter : 600 μm) and scan speed were varied between 700 and 1000 W, 4.25 and 21.2 mm/s, and the layer thickness and hatch distance were kept constant to 0.2 and 0.3 mm, respectively. The printing chamber was filled with argon (99.9%, Matheson) to prevent oxidation of the Cu and D powders. The printing parameters are summarized in Figure 5.2-1.

Figure 5 . 2 - 1 (

 521 Figure 5.2-1 (a) Table of printing parameters and (b) schematic illustration of the

5 . 2 - 2 .

 522 microscope (SEM) and a high-resolution transmission electron microscope (HRTEM, FEI Tecnai Osiris ™ , 200 kV). The periodic photothermal radiometry (PPTR) 238 was used to measure the thermal diffusivity (TD) of a printed Cu/D composite material as schematically illustrated in Figure The setup is equipped with an Nd:YAG laser (Coherent MATRIX Q-switched diode-pumped laser, λ = 1064 nm) that delivers a 40 ns pulse duration each period TR with a maximum power of 10 W. The laser beam (diameter 0.55 mm) was directly focused on the sample; an infrared (IR) detector (Teledyne Judson Technologies J15D12) was used to monitor the temperature change at the rear face of the sample.

Figure 5 .

 5 Figure 5.2-2 (a) Schematic illustration of the laser-flash setup for TC measurement and

Figure 5

 5 recoating and a remelting strategy were employed at scan speeds and hatch distances ranging from 200 to 450 mm/s and 0.6 to 0.12 mm. The remelting strategy involves shinning the laser twice on the same layer, as shown in Figure 5.2-3 (b). The recoating implies repassing the recoater on the same layer and shining the laser a second time, as illustrated in Figure 5.2-3 (c).The structures designed using SolidWork® were uploaded to the Magics® software for parameters assignment and print job creations. Then, the parts were printed on a stainless-steel building plate (stainless steel, 123 × 123 × 25 mm 3 ). The SLM was monitored using a layer control system (LCS) installed in SLM®125 printer.

Figure 5 . 2 - 3

 523 Figure 5.2-3 Schematic illustration of (a) single melting, (b) remelting, and (c) recoating

4 ,

 4 a Titanium (Ti)-based coating on carbon fibers, via a molten salt process, has demonstrated an enhanced wettability of molten Cu on a graded TiO2-TiC layer compared to a pure TiC layer. The same process has been applied to the D particles. 240 Coated D particles were then mixed with Cu powder Figure 5.1-1 (b)). The printing parameters were explored to form a dense Cu/D composite material. It was observed that scan speeds between 8.7 and 21.7 mm/s and laser powers between 700 and 1000 W, mainly resulted in parts with large numbers of printing defects on the surfaces (i.e., pores, balling) (Figure 5.3-1 (a) and (d)). The poor surface finish has arisen because the Cu powder was insufficiently melted which inhibits its ability to fuse completely. 126 Nevertheless, a narrow processing window was identified for obtaining parts with minimized printing defects: a scan speed of 12.7 mm/s and laser powers ranging from 900 to 925 W (Figure 5.3-1 (b) and (d)). At a low scan speed (4.3 mm/s) and high laser powers (800 to 1000 W), the printed parts exhibit a spherical shape, not a designed cubic shape (Figure 5.3-1 (c) and (d)). Slow scan speeds and high laser powers enlarge the molten pool size, which increases the powder track height and width. As a result, the original cube deforms into a spherical shape. 241

Figure 5 . 3 - 1

 531 Figure 5.3-1 SEM micrographs of Cu/D composite surfaces printed at an energy density of (a) 708, (b) 1180, and (c) 1764 J/mm 3 ; (d) qualitative quality of Cu/D vs. printing

Figure 5 . 3 - 2 (

 532 c) shows a cross-sectional micrograph of the interface between Cu and D, where the coated-D was previously embedded in the Cu matrix (Figure 5.3-2 (a) and (b)). As shown in the SEM micrographs, an intimate contact between the Cu matrix and the coated-D particle was found. The HRTEM micrographs at the Cu-TiO2 and TiO2-TiC interfaces demonstrate a smooth transition between layers with no pores, confirming a chemical bonding (Figure 5.3-2 (d) and (e)). Additionally, HRTEM of the Cu, TiO2, and TiC layers in Figure 5.3-2 (f)-(h) shows the crystalline structures of the (111) Cu, (101) TiO2 and (111) TiC planes with interlayer distances of 0.21, 0.35, and 0.25 nm, respectively. 243,244

Figure 5 . 3 - 2 (

 532 Figure 5.3-2 (a)-(b) SEM micrographs of the Cu/D surface. TEM micrographs of (c)

Figure 5

 5

Figure 5 . 3 - 3

 533 Figure 5.3-3 SEM micrographs of D particles at different stages of (a) as received, (b)

Figure 5 . 3 - 3 (

 533 d) and (e), the processed D particles still exhibit a sharp D peak, suggesting no graphitization occurred during the hightemperature coating and printing processes. The D peak position and FWHM values retrieved from the Raman spectra show no obvious peak broadening, confirming the multistep processing does not degrade the crystallinity of the D particles.

  of Cu/D via a single laser melting strategy As shown previously, the manufacturing of Cu/coated-D composites via LDED exhibits a chemical bonding between the Cu matrix, the (Ti)-based interphase, and the D particles leading to the fabrication of highly thermally conductive MMCs. 247 Nevertheless, as mentioned in Chapter 2, the LDED is a technology used to print large objects or repair parts. Thus, the fabrication of Cu/D composite material is then studied via SLM. Note that, a similar coating as shown in the previous part was realized here on D particles, as presented in Figure 5.1-1 (d).

Figure 5 . 4 - 1 (

 541 a) to (c) show the SEM micrographs of printed Cu/D composites with a low, moderate, and high energy densities of 333, 533, and 1333 J/mm 3 , respectively. The surface of all samples appears porous, with pore/cavity sizes ranging from 100 to 500 μm. The density of the samples increases from 67% to 90% when the energy density increases (Figure5.4-1 (d)). Previous reports demonstrated that an higher energy density (i.e., lower scan speed and hatch distance)

Figure 5 . 4 - 1

 541 Figure 5.4-1 SEM micrographs of Cu/D surface printed with energy densities of (a) 333, (b) 533, and (c) 1111 J/mm 3 and (d) Cu/D density vs. energy density.

Figure 5 . 4 - 2

 542 Figure 5.4-2 Optical photograph of the laser shining on (a) Cu and (b) Cu/D powder

Figure 5 . 4 - 3

 543 Figure 5.4-3 Schematic illustration of spatter ejection in SLM.

Figure 5 . 4 - 4 (

 544 a) shows the pictures of the powder bed after the coating of a new layer and at different stages of the printing for a single melting (i.e.,layer 10, 60, 110, and final part). As can be seen, at the beginning of the print (layer 10 ~ 300 μm), the powder bed is flat, uniform, without defect. Next, when the number of layers increases (layer 60 ~ 1200 μm), white and large particles appear on the surface. Later, they disappear (layer 110 ~ 3300 μm), and the printed part emerges at the powder bed surface, witnessing its high surface roughness. The fading away of the large particles after 110 layers is due to their incorporation into the printed part, causing the roughness of the parts and their absorption inside the bed. Sonny Ly et al. reported that the size of the recoil-induced particle ejection

Figure 5 . 2 - 3 (

 523 photographs for the remelting strategy and shows that the large spatter particles have

Figure 5 . 4 - 4

 544 Figure 5.4-4 Pictures of the powder bed after the coating of a new layer using the (a)

Figure 5

 5

Figure 5 . 4 - 5

 545 Figure 5.4-5 Photographs of Cu/D composite materials printed with scan speeds and

Figure 5 . 4 - 6

 546 Figure 5.4-6 2D map of the part quality as functions of the laser scan speed and hatch

  microstructures. All samples were manufactured with a constant scan speed and hatch of 250 mm/s and 0.1 mm, respectively (Figure 5.4-7 (a)). Figure 5.4-7 (b) shows the bottom, top cross-section, and surface of the Cu/D composite printed with a single melting strategy.

Figure 5 . 4 - 7 (

 547 c) shows the microstructure of the sample printed using the recoating strategy. The bottom part of the sample is denser compared to the single-melting approach, but the inclusion remains.

Figure 5 . 4 - 7 (

 547 d) displays the microstructure of the sample printed using the remelting strategy. The bottom of the sample appears to be dense, with only a few inclusions. The bottom of the sample with no defect leads to the formation of a dense composite during the layer-upon-In this part, the effects of the single melting, recoating, and remelting printing strategies are systematically compared in terms of powder bed quality, part density, and microstructure of the Cu/D composites laser 3D print. It is shown that the presence of coated-D influences the molten pool dynamics during the SLM, generating a large amount of spatter ejections. The spatters are assumed to be originated from the recoil-induce spatter due to the high metal vapor above the pool when D are introduced in the powder mixtures.Later it is shown that the spatter perturbed the powder bed homogeneity and resulted in the formation of porous composite materials. Besides, the low laser absorption of Cu causes significant back reflection, and limits the laser scan speed to a minimum of 200 mm/s. Later, to obtain dense composite material, both the recoating and remelting strategies are introduced during the SLM. It is shown the only the remelting approach reduces the spatter ejection by suppressing the printing defects and thus dense Cu/D composites (~96%) are manufactured with scan speeds of 200, 250, 300 and 350 mm/s and hatch distances of 0.12, 0.10, 0.08 and 0.06 mm, respectively. Finally, complex 3D structures are printed for thermal management applications.Honeycombs of Cu/D with different fillings and D-structures are successfully manufactured using the remelting strategy with a scan speed of 250 mm/s and hatch of 0.10 mm. The introduction of the remelting in the SLM has promising prospective to laser 3D print a large variety of MMCs and thus remodel their applications via the development of intricated designs.

  of the UV-laser-assisted combustion D CVD experiment setup is shown in Figure6.2-1. A combustion torch with a 1.5 mm orifice tip was used to produce the flames. The precursor gases were a mixture of ethylene (C2H4, 99.999%), acetylene (C2H2, 99.6%), and oxygen (O2, 99.996%) with a volume ratio of 1:1:2. A krypton fluoride (KrF) excimer laser (Lambda Physik ® , COMPex ® 205, 248 nm, 20 ns) and an argon fluoride (ArF) excimer laser (Lambda Physik ® , COMPex ® 205, 193 nm, 15 ns) were directed perpendicularly through the combustion flames and parallel to the substrate to excite the combustion species. A UV convex lens with a focal length of 25 mm was used to focus the laser beam from its original size of ~ 20×10 mm 2 to ~ 2×1 mm 2 , which was sufficient to cover the whole primary flame. A power meter was placed in the UV laser path next to the combustion flame to measure the laser power. The laser fluence was calculated by dividing the laser pulse energy by the beam spot size. D growth was performed with UV laser irradiation of the combustion flame by tuning the laser fluence from 600 to 1400 mJ/cm 2 and from 10 to 25 mJ/cm 2 for the 248 and 193 nm lasers, respectively. A tungsten carbide (WC) substrate with a Co composition of 6% and a dimension of 12.5×12.5×1.6 mm 3 was placed on a water-cooled brass plate that was mounted on an X-Y-Z stage. The distance between the substrate surface and the primary flame tip was kept around 0.5 mm in all experiments. The substrate temperature during the deposition was maintained at 770 ~ 780 ℃ and monitored by a noncontact pyrometer (OS3752, Omega Engineering, Inc.). For the investigation of the growth rate and film quality, the deposition time was varied to obtain a similar film thickness, ~10 mm, for comparison purposes. For the D nucleation study, the deposition time was 10 min.

Figure 6 . 2 - 1 6 . 2 - 2 .

 621622 Figure 6.2-1 Schematic illustration of the diamond-assisted UV laser setup

Figure 6 . 2 - 2

 622 Figure 6.2-2 Schematic diagram of the experimental system used for optical emission

Figure 6 .

 6 3-1 (a) illustrates the acquisition time and sequences between the KrF UV laser and the intensified charge-coupled device (ICCD) gate of the spectrometer. The emission peaks from three mains species were detected in the spectra (Figure6.3-1 (b) C2: X' 3 Pu ← A 3 Pg (Dv = -1, 0, 1, 2), 2) CH: A 2 Δ ← X 2 P (Dv = 0), and 3) OH: A 2 Σ + ← X 2 P (Dv = 0). All emission peaks of the species increase as the laser fluence increases.

Figure 6 . 3 - 1 . 6 . 3 - 2 . 6

 6316326 Figure 6.3-1. (a) The time relationship between the synchronized ICCD gate of the

Figure 6 . 3 - 3 (

 633 Figure 6.3-3 (b). Significant enhancement of the peak intensity at 512.9 nm was observed

Figure 6 . 3 - 3 (

 633 c) shows CH was excited through the X 2 P (v"=0) → B 2 S -(v'=0) transition at 388.9 nm. Because the B 2 S -state and the A 2 D state crossed each other, strong collisional redistribution between these two states took place.The fluorescence from the A 2 D (v'=0) → X 2 P (v"=0) transition was observed around 431.4 nm. Significant enhancement of the peak intensity at 431.4 nm was observed in the LIF spectra, confirming the effective excitations of the ground-state CH species by the OPO laser probing.

Figure 6 . 3 - 3 (

 633 d) presents the LIF detection of OH was carried out by excitation of the X 2 P (v"=0) → A 2 S + (v'=0) transition at 307.8 nm. Detection of fluorescence around 343-348 nm (A 2 S + (v'=0) → X 2 P (v"=1)) was observed. Quantitative analysis of the LIF signals was impeded by quenching the fluorescence due to inelastic collision of the excited molecules. Certainly, collisional quenching cannot be neglected in flames at the atmospheric pressure. This quenching depends strongly on the pressure, temperature, and collision partners, which are different in different regions of the flame. The LIF results were, therefore, interpreted in a more qualitative way. The LIF signals at 512.9, 431.4, and 347.2 nm in Figure 6.3-3 (b)-(d) were integrated to represent the abundance of the ground-state C2, CH, and OH species in the flame, respectively.

Figure 6 . 3 - 3

 633 Figure 6.3-3 The time relationship among the synchronized ICCD gating of the

Figure 6 . 3 - 4

 634 Figure 6.3-4 Integrated intensities of (a) OES peaks and (c) LIF signals of C2, CH, and

Figure 6 . 3 - 5 (

 635 b) shows the Boltzmann plots of ln(Iλ 4 /SJ'J'') versus Ej'/kB, in which the slopes correspond to 1/Tr.

Figure 6 . 6 . 3 - 6 ,

 6636 Figure 6.3-5 (a) High-resolution optical emission spectra of the rotational R-branch

Figure 6 . 3 - 6

 636 Figure 6.3-6 Flame temperature as a function of the KrF excimer laser fluence.

Figure 6 . 3 - 7

 637 Figure 6.3-7 Diagram of transition processes of reactive radicles, C2, CH and OH, under

Figure 6 . 3 - 8

 638 Figure 6.3-8 (a). The diamond film deposited without KrF excimer laser irradiation

Figure 6 .

 6 Figure 6.3-8 (a) SEM micrographs of surface and cross-sectional morphologies of the D

Figure 6 .

 6 Figure 6.3-9 (a) Schematic illustration of in situ field-enhanced thermionic current

Figure 6 . 3 -

 63 Figure 6.3-11 (a) shows a TEM micrograph of the cross-section of a 389 nm thick D film on WC prepared with KrF excimer laser irradiation at 1.4 J/cm 2 . A thin platinum (Pt)

Figure 6 . 3 -

 63 Figure 6.3-11 (a) A TEM micrograph of a 389 nm thick D film on a WC substrate prepared with KrF excimer laser irradiation at 1.4 J/cm 2 . HRTEM micrograph of (b) the

Figure 6 . 3 -

 63 Figure 6.3-12 (a) A TEM micrograph of a D film sample prepared without KrF excimer

6. 4 6 . 4 . 1 Figure 6 . 4 - 1

 4641641 Effect of UV laser wavelength (KrF, λ = 248 nm vs. ArF λ = 193 nm) on the diamond growth Effect of UV wavelength on D growth rate and quality To identify how UV laser irradiation at different wavelengths affects D growth, D films were prepared under ArF and KrF excimer laser irradiations at 193 and 248 nm, respectively. shows the SEM micrographs of the surface morphologies of the D films after 1 h of deposition.

Figure 6 . 4 - 1 6 . 4 - 2 ,

 641642 Figure 6.4-1 SEM micrographs of surface morphologies of the D films prepared for 1 h

Figure 6 . 4 - 2

 642 Figure 6.4-2 Thickness profiles of D films prepared at different laser fluences at a UV

Figure 6 . 4 - 3 6

 6436 Figure 6.4-3 Growth rates of the D films prepared for 1 h deposition under 193 and 248

Figure 6 . 4 - 4

 644 Figure 6.4-4 Raman spectra of the D films prepared with 193 or 248 nm laser

Figure 6 . 4 - 5 D

 645 Figure 6.4-5 D quality factors of the diamond films plotted as functions of the laser

Figure 6 . 4 - 5

 645 Figure 6.4-5 shows the quality factors plotted as functions of the laser fluence for the 193

Figure 6 . 4 - 7 ( 6 . 4 . 2 Figure 6 . 3 - 1 Figure 6 . 4 - 8 )

 647642631648 Figure 6.4-7 (a) nucleation times as functions of UV laser wavelengths and fluences,

Figure 6 . 4 - 8 6 . 4 - 9 .

 648649 Figure 6.4-8 OES spectra of the combustion flame under 193 or 248 nm laser irradiations

Figure 6 . 4 - 9

 649 Figure 6.4-9 Flame images with and without the filters under different laser irradiation

6

 6 determine how the species were generated by UV laser irradiations, the flame temperatures were estimated for both laser wavelengths at different fluences. Due to a strong coupling between the translational and rotational energy states of the molecules, the flame temperature was approximated by the rotational temperature derived from high-resolution rotational line emission intensities of CH using the Boltzmann plot (Figure

Figure 6 . 4 -Figure 6 . 4 - 12 28 ) 29 )Figure 6 . 4 - 13

 64641228296413 Figure 6.4-10 High-resolution optical emission spectra of the rotational R-branch of the

  photodissociation of C2H4 has also been well characterized.311 Ethylene molecules can be

  The thermal properties and heat dissipation performance of the Cu/D structures printed should be the focus of future work. Next, only a low concentration of D particles was introduced in the Cu matrix (i.e., 5 vol.%). It is thus essential to increase the D concentration gradually and to investigate the effect of D concentration on the printing parameters and thermal characteristics. This study could result in the ability to anticipate the printing parameters and thus easily produce tailored TC and CTE of Cu/D materials.

Table 1 .

 1 Physical

[START_REF] Mortimer | The wetting of carbon by copper and copper alloys[END_REF][START_REF] Veillère | Relationship between interphase chemistry and mechanical properties at the scale of micron in Cu-Cr/CF composite[END_REF][START_REF] Chung | Effect of Titanium Addition on the Thermal Properties of Diamond/Cu-Ti Composites Fabricated by Pressureless Liquid-Phase Sintering Technique[END_REF]

...................................................................................................................181 properties of the printed and the commercial (COM) columnar heat sinks and bulk Cu made by powder metallurgy route. ................................................................

55 Table 4.3-1 XPS quantitative analyses using peak integration

  . ........................................77

  149,150 Copper (Cu) reinforced carbon fibers (CFs), or diamond (D) are excellent examples of MMCs with weak interfacial bonding. These MMCs are of particular interest in thermal management, as next-generation heat sink materials.

151 Solid Cu alone is limited by fatigue failure due to stresses of thermal cyclicals. Adding carbon to reinforce Cu offers an advantage of combining the low CTE and high TC carbon (C) with the excellent TC of Cu (400 W/m.K). Nevertheless, a simple Cu/C composite suffers from

  [272][273][274][275][276]

	variant of water, HOD, to influence outcome through selective bond cleavage with
	wavelength-matched infrared (IR) laser irradiation. 277 Ding et al. showed that ultraviolet
	(UV) photolysis of glyoxal, (CHO)2, proceeded through different pathways depending on
	the photolysis wavelength, subsequently leading to different products. 278 Laser-assisted
	CVD (LCVD) opens up new possibilities in thin-film deposition and enables one to study
	new reaction pathways.

(Co) (dimensions: 12.5×12.5×1.6 mm 3 ). Samples were placed on a water-cooled brass stage 10 mm beneath the tip of the flame for 60 s. The temperature was evaluated by a noncontact pyrometer (OS3752, Omega Engineering, Inc.) and measured to be 1200 ºC.

Characterization methods

Morphological and microstructural characterization was performed using an SEM (VEGA II SBH, TESCAN). Crystal structures and chemical phases were identified by Xray diffraction (XRD) (PANalytical X'pert PRO MPD, λ = 1.5405 Å). The stoichiometry of TiCx was estimated by calculating the lattice parameters from the XRD diffractograms via the Le-Bail method 185 using the FullProf Suite. Auger electron spectroscopy (AES) (Auger VG Scientific Microlab 310F) and X-ray photoelectron spectroscopy (XPS, Thermo Electron K-alpha photoelectron spectrometer, Al Kα excitation) were used to determine the elemental and chemical compositions. The AES sputtering time was converted to a depth profile by measuring the layer removal depth using SEM. Grain morphology, orientation, and size were evaluated using electron backscatter diffraction (EBSD). The oxidation state and crystallinity of the coated CFs after being subjected to the oxy-acetylene flame were investigated using a micro-Raman spectrometer (inVia ™ , Renishaw). An Ar + ion laser with a wavelength of 514.5 nm and a power of 50 mW was used as the excitation source. The laser beam was focused using a 50× objective lens to approximately 5 µm. layer process. These observations confirm the advantages of using the remelting printing strategy to manufacture dense micro-reinforced MMCs.

3D printing of complex Cu/D structures

Cu/D composite materials are commonly referred as the next generation of thermal management materials to be utilized as heat sinks and base plates (Figure 5.4-8 (a)). Heat sinks require high surface areas to enhance the heat dissipation, hence intricated designs.

However, these intricated designs are difficult to achieve on D-based materials. [START_REF] Neugebauer | Additive manufacturing boosts efficiency of heat transfer components[END_REF] On the other hand, Cu has a high density (8.93 g/cm 3 ) withstand their applications compared to Al (2.7 g/cm 3 ), which is widely used in thermal applications. Nevertheless, SLM has the ability to fabricate efficient and lightweight Cu/D heat dissipators by forming intricate and porous designs. Therefore, different complex structures were printed, using the optimized parameters developed in the previous parts (i.e., remelting s = 250 mm/s, h = 0.1 mm), which have prospective uses in heat dissipation. First, a diamond lattice structure was printed, as shown in Figure 5.4-8 (b). The SEM micrographs show the sophisticated and thin features (400 μm) of the D-structures printed. The thermal performance of a Dstructure made of Al alloys was previously reported by M. Wong et al.. [START_REF] Wong | Pressure Loss and Heat Transfer Through Heat Sinks Produced by Selective Laser Melting[END_REF] They showed that the heat dissipation of a D-structure is 60% more efficient than an array heat sink while being lightweight and strong. 258,259 Hence, D-structure made of Cu/D composite is likely to show excellent heat dissipation performance when used as a heat sink. In addition, three honeycomb structures were additively manufactured with filling percentages ranging from 40 to 60% (Figure 5.4-8 (c)) Honeycombs are a well-established structure as effective and lightweight heat dissipators. 260,261 Combining this shape with the excellent thermal properties of Cu/D materials could improve the thermal performance of base plates by increasing the surface areas and lowering their weights. In summary, the introduction of the SLM method to manufacture Cu/D composite materials has immense potential to produce thermally efficient designs and surpass the thermal properties of available heat dissipators components. Three filters were placed in the optical path to measure the 2D emission images created by the emission band from each excited species: diatomic carbon (C2), methylidyne radical (CH), and hydroxyl radical (OH). Upon inspection of the images, it is clear that UV laser irradiation effectively induces generation of all three excited species, following the same trend in the spectra. It is noted that the increment of OH is concentrated in the center of the flame, while C2 and CH are more significant at the flame edges. A limitation of OES is that it only probes excited-state species, which are usually not sufficient in the overall chemistry because their number densities are several orders of magnitude smaller than those at the ground states. LIF is often the technique of choice to selectively and unambiguously detect ground-state species by its excitation spectrum. The sum of the integrated intensities of the non-D carbon bands, respectively 292 . The D quality factor exhibites a nearly linear increase with respect to the laser fluence, reaching a value of 97.4% at a laser fluence of 1.4 J/cm 2 , 4.2% higher than the film prepared without KrF excimer laser irradiation (Figure 6 energy leading to direct photodissociation of both C2H2 and C2H4 via either one-photon or two-photon processes. Photolysis rate, the number of times a photodissociation occurs per photon absorbed, is highly wavelength dependent. The shorter the wavelength of the radiation is , the high photolysis rate would be obtained. Although C2H4 has a C-H bond energy which is around 0.2 eV below the 248 nm photon energy, the photolysis rate is much lower than that at 193 nm, of which the photon energy is around 1.5 eV higher than the C-H bond energy. Compared with 193 nm laser irradiation, direct photodissociation of either C2H2 or C2H4 by absorbing the 248 nm photon is difficult and must proceed through a two-photon process, which requires high laser intensity. This explains why, although the output fluence of the 193 nm laser was one order of magnitude lower than that of the 248 nm laser, the 193 nm laser irradiation of the combustion flame led to a more pronounced enhancement in diamond growth and a more significant increase in the generation of key radicals.

Conclusions

The effects on diamond growth through UV laser irradiation of the D-forming combustion flame at different laser wavelengths is studied. The 193 nm laser irradiation is found to be more efficient in enhancing D growth and modifying the flame chemistry than the 248 nm laser irradiation. Similar D growth enhancement is obtained with laser irradiations at both 193 and 248 nm, although the laser fluence of the 193 nm laser is one order of magnitude lower than that of the 248 nm laser. Investigation of the lateral grain size evolution and the D nucleation process suggest that the secondary nucleation is significantly suppressed by the UV irradiation to the flame at 193 nm. A higher increment rate of reactive species under the 193 nm laser irradiation is identified by analyzing the flame chemistry using OES. The more pronounced effects of the 193 nm irradiation can be attributed to the high photonic energy which leads to the direct photodissociation of hydrocarbon precursors.

Summary and Outlook

Summary

The objectives of the research presented herein were to additively manufacture intricate Cu/D composite materials designed for heat sinks and achieve fast, high-quality deposition of D films to be used as heat spreaders. Several challenges were addressed to fulfill the above objectives, such as the difficulty of 3D laser printing of Cu/D composite materials, the lack of chemical affinity between both Cu and D components, and the low D film quality obtained using combustion flame.

The research first focused on the 3D laser printing of pure Cu. The SLS of dense, smooth, and crack-free Cu structures was demonstrated using a low-power green laser (5 W, λ = 532 nm) by compacting the powder. Cold pressing the Cu powder enhanced the sintering process and limited the formation of printing defects during SLS. Then, SLM of pure Cu was investigated using a NIR laser (400 W, λ = 1070 nm). After optimizing the printing parameters, it was demonstrated that laser power, scan speed, hatch distance, and layer thicknesses of 400 W, 400 mm/s, 0.12 mm, and 0.03 mm, respectively, led to the manufacturing of dense and crack-free Cu parts (95%). The measured properties of a simple printed heat sink were equivalent to a commercial one having the same geometry with an electrical and thermal conductivity of 98% and 100% IACS and 368 and 378 W/m.K, respectively. The advantages of SLM over conventional manufacturing methods