5

In the last 20 years, the key importance of energy and environmental challenges has paved its way to vehicle consumers. They are becoming more and more aware of environmental issues and are looking for cleaner vehicles. In parallel with the customer's desire, accelerated regulatory constraints on CO2 emissions and pollutants (such as CO, NO2) for new vehicles came into force in order to reduce greenhouse gas emissions. In this context, light-weight and electrical vehicles-related materials are one of the most important trends for the development of materials in the transportation sector. This PhD work aimed at exploring novel metal matrix nanocomposites fabricated by ultrasound-assisted casting techniques to address automotive-related issues: mechanical enhancement and thermal management, with specific emphasis on the processing parameters for the dispersion ability and the interaction between nanosized-inclusion and metal matrix.

Aluminum and its alloys have been largely used in many industrial fields because of their excellent properties, such as lightweight, low cost, good thermal properties and high ductility. But the poor stiffness, strength, and wear properties still limit their applications as vehicle materials. Metal matrix composites are a class of metallic materials containing reinforcements and that can be traced back to the 1970s. Nanoparticles could be applied in a similar manner to enhance the mechanical properties of aluminum-based matrices by the design of metal matrix nanocomposites, which will be discussed in chapter 1. In the same chapter, the mechanisms behind the improved properties of nanocomposites will also be described. Then the techniques used to fabricate metal matrix nanocomposites will be further discussed. Chapter 2 will present the synthesis and characterization of building blocks for the preparation of nanocomposites for thermal management: core-shell Bi@SiO2 nanoparticles obtained by sol-gel synthesis. Our efforts for decreasing the aggregation state of Bi nanoparticles in silica shell will be presented, including ligand exchange at the surface of Bi nanoparticles. These particles will be used as building blocks of composites in chapter 4.

The incorporation of stiffer and harder reinforcement materials leads to aluminum metal matrix composites combining the properties of both the matrix and the reinforcement phase.

Considerable efforts have been deployed to incorporate commercial nanoparticles into metal matrices. Many recent chemical synthesis methods have widened the range of accessible nanoparticles, especially for boron-based phases, which are interesting for vehicle applications due to their light weight and hardness. Their application as metal matrix inclusions is studied and will be discussed in chapter 3. This chapter will present the fabrication and characterization of HfB2-Al and SiC-Al nanocomposites. It will present in detail the original set-up we have designed, based on ultrasound-assisted casting, in order to increase the dispersion of nanoparticles in aluminum-based matrices. In the same chapter, the microstructure evolution with different experimental parameters is determined and the mechanical properties of these composites are evaluated by microindentation and nanoindentation tests.

While electrical vehicles are promising systems and have increased market share since several years, the research on the thermal management of battery packs is still in need. Among classic cooling methods, phase change material is a passive strategy relying on heat absorption through melting. Several recent developments focused on metal matrix nanocomposites with phase change materials as inclusions, which are of great interest due to their high thermal conductivity, high mechanical properties and wide working temperature range. Conventional strategies including powder metallurgy and bottom-up fabrication require high fabrication costs. At the operational level, we investigated the liquid phase method: metallic nanoparticles were encapsulated in a shell of SiO2 to prevent destruction of the particles and macroscopic phase separation. Subsequently, these nanoparticles are incorporated in the molten metal matrix by the ultrasound-assisted process developed in chapter 3. This whole part will be presented in chapter 4, which encompasses the synthesis of phase change nanocomposites, the study of the reactivity occurring at the nanoparticles-matrix interface, and the assessment of the heat absorption ability of the resulting composites.

Overall, this work opens new perspectives for the application of ultrasound-assisted casting to the fabrication of mechanically reinforced and phase change metallic materials that could be potentially applied to fabricate vehicle components. We have also identified specific difficulties in reaching high dispersion of boron-based nanoparticles in aluminum matrices, thus identified future tracks to explore in order to achieve metal matrix nanocomposites with high homogeneity at the nanoscale for a wide range of reinforcement nanomaterials.

Chapter 1 Bibliography and state-of-the-art

This first chapter focuses on bibliographic aspects. First, the notion of metal matrix nanocomposites and their actual/potential applications in the automotive industry are presented. Then, the interest shifts toward two potential applications, mechanical reinforcement and heat management, with overviews of the existing techniques and solutions.

Finally, the fabrication methods for nanocomposites are briefly presented.

Context

Composite materials consist of two or more constituents with different physical and/or chemical features. They can exhibit properties superior to each of the components and to their sum, thanks to so-called synergistic effects. 1 The major constituent is regarded as the matrix phase, and the remaining is called the reinforcement phase(s). 2 Often, the reinforcement phase is stronger, harder and stiffer than matrix phase, enabling improvement of the properties of the matrix. Commonly, composite materials are divided into three main types according to the matrix phase: polymer matrix composites, 3,4 metal matrix composites (MMCs) 5,6 and ceramic matrix composites. 7,8 Especially, MMCs have attracted lots of attention due to their superior properties and wide applications range from aerospace, to automobile industry and to defense.

Metal matrix composites

Metal matrix composites (MMCs) are a class of metallic materials containing reinforcement phase(s) that can be traced back to the 1970s. Researchers managed to improve properties of metals/alloys with ceramic reinforcements by stir casting, rheocasting and powder metallurgy. 9,10 For MMCs, three main types of reinforcement materials can be distinguished: continuous fiber networks and discontinuous reinforcement materials as particles or whiskers 11 (Figure 1.1). Compared to continuous reinforcement, the discontinuous one presents advantages of lower cost and isotropic properties. Generally, there are two main types of reinforcement phases used for MMCs: (a) ceramic reinforcements, such as Al2O3, 12,13 SiC, 14,15 TiC, 16 B4C, 17 and (b) carbon-based materials like graphite, fly ash, etc. To fabricate MMCs, aluminum, 2,18 magnesium, 19,20 titanium 21 and copper 22 are widely used as matrix phases. 11 As mentioned above, reinforcement with metallic or ceramic phases can improve hardness of the composite versus the sole matrix, enabling the material to withstand higher load but also providing additional specific properties. The incorporated reinforcement objects could be microsized or nanosized. Micron-scale reinforcement in the metal matrix has been well studied and demonstrated to provide high hardness, elastic modulus, and strength. 23 Reducing the size of the inclusions to the nanoscale brings additional advantages: reduced load (5 to 20% for macroscale reinforcements vs. lower than 5% for nanoscale ones) and increased strength by Orowan hardening and pinning of grain boundaries. 24 Hence, metal matrix nanocomposites (MMNCs) possess in general properties superior to those of composites with microscale reinforcement. Nanoscaled reinforcement materials for MMNCs include graphene, carbon nanotubes and nanofibers, and nanoparticles (such as Al2O3, SiO2, SiC, TiO2). 5,6

Figure. 1.1 Three types of reinforcement in MMCs

How can aluminum matrix nanocomposites improve sustainability in the automotive industry?

A new form of mobility represented by the ACER concept (Autonomous, Connectivity, Electrification and Ridesharing) has shifted dramatically the market in the past few years.

Based on the latest legislation and current trends dealing with clean vehicles and carbon neutrality, automotive industries have integrated sustainability into their strategy. In this context, the research of light-weight, tough, and functional materials has a growing significance.

The application of nanocomposites for the design of automotive components reflects the search for improved production rate and thermal stability, weight reduction, and less wear parts. All these features contribute to a reduction of CO2 emissions and of environmental impact. Most of them also enable an increase of the cruising time, which is of great importance in the case of electric vehicles. 25 Light (density < 2.7 g•cm -3 ) and tough (tensile strength > 350 MPa) aluminum MMNCs are especially interesting for automobile applications. Aluminum and its alloys are already largely used in many industrial fields because of their light weight, low cost, good thermal stability and high ductility. But their poor stiffness, strength, and wear properties still limit their applications. 2 These drawbacks can be circumvented by the design of aluminum metal matrix composites with the addition of reinforcement materials stiffer and harder than aluminum and its alloys.

Apart from traditional mechanical reinforcement, metal matrix nanocomposites can also acquire properties specific to the inclusion nanoparticles. A typical application deals with thermal management of batteries. 26 The combination of phase change inclusion particles and aluminum matrix leads to bifunctional materials with heat management properties and good thermal dissipation. 27 These properties will be detailed in the next chapters.

Mechanical property improvement

Enhanced mechanical properties is one of the main pursuits for MMNCs fabrications. Below, we present possible strengthening mechanisms and strategies to improve mechanical properties.

Strengthening mechanisms

Most of the work on metal matrix nanocomposites focuses on the improvement of mechanical properties: hardness, rigidity and wear resistance. 23 As the major component, the metallic matrix provides ductility. Its properties depend on the polycrystallinity of its microstructure. Indeed, grain boundaries strengthen the material by acting as barriers to the dislocation movement.

The enlisted strengthening mechanisms for MMNCs encompass load transfer effect, Hall-Petch strengthening, Orowan strengthening, coefficient of thermal expansion (CTE) and elastic modulus (EM) mismatch. 24 Load transfer effect corresponds to direct strengthening by transferring of the load from the matrix to the reinforcement nano-objects. The reinforcement material is harder than the matrix, thus contributing to strengthening of the overall material. Nardone and Prewo 28 proposed a modified Shear lag model to explain the effect of load transfer on mechanical properties. Since the volume fraction of the reinforcement particles is small, load transfer is small as well. For isotropic particles, the model can be simplified to yield the equation 29 below:

∆s 𝐿𝑇 = 1 2 𝑣 𝑃 𝑠 𝑚 (1)
Where 𝑣 𝑃 and 𝑠 𝑚 are the particle volume fraction and the yield strength of the un-reinforced matrix, respectively.

Orowan strengthening. This effect is based on the interaction between nanoparticles and dislocations. It can be ignored for particles with a diameter larger than 1 µm. 30 Dislocations tend to bend around particles and to form dislocation loops with dislocations on the other side of the particles. These dislocation loops are then pinned onto the nanoparticles, which leads to enhanced strength. At same volume fraction, smaller particles contribute to the higher enhancement of final strength. This effect can be expressed as the equation 31,32 below:

∆𝑠 𝑂𝑅 = 0.13𝑏𝐺 𝑚 𝑑 𝑝 [( 1 2𝑣 𝑝 ) 1 3 -1] ln 𝑑 𝑝 2𝑏 (2)
where 𝑏, 𝐺 𝑚 , 𝑑 𝑝 and 𝑣 𝑝 are the Burgers vector, shear modulus, particle average diameter and particle volume fraction, respectively. For example, Cao et al. [START_REF] Cao | Study on tensile properties and microstructure of cast AZ91D/AlN nanocomposites[END_REF] applied ultrasound-assisted casting to fabricate AlN reinforced magnesium alloy AZ91D that exhibited a yield strength of 144 MPa compared to 100 MPa for the nonreinforced alloy. The author ascribed this 44% increase to the Orowan strengthening and the effect of the particles decreasing (or refining) the matrix grain size (see below), contributing to 35.8 MPa and 8.2 MPa, respectively.

Hall-Petch strengthening. The improvement of strength is also related to decreased grain size since grain boundaries hinder dislocation movements. 24 The relationship between grain size (𝑑) and yield strength (∆𝑠 𝐻-𝑃 ) is expressed through the Hall-Petch equation: [START_REF] Hull | Introduction to dislocations[END_REF] ∆𝑠 𝐻-𝑃 = 𝑘 𝑦 √𝑑 (3) where 𝑘 𝑦 is the strengthening coefficient. From equation (3), we can clearly see that the yield strength of nanocomposite is inversely correlated to the grain size. This is true for particle sizes above ca. 10 nm. Below, the inverse Hall Petch effect prevails, leading to an opposite dependence versus the grain size. [START_REF] Sanaty-Zadeh | Comparison between current models for the strength of particulate-reinforced metal matrix nanocomposites with emphasis on consideration of Hall-Petch effect[END_REF] Hence, an optimal grain diameter for Hall Petch hardening is ca. 10 nm.

The addition of particles results in a decrease of the grain size of the matrix through localization of the particles at the grain boundaries, which retards or stops matrix grain growth. 24 The Zener equation [START_REF] Sanaty-Zadeh | Comparison between current models for the strength of particulate-reinforced metal matrix nanocomposites with emphasis on consideration of Hall-Petch effect[END_REF] shows the relationship between grain size (𝑑 𝑚 ), particle volume fraction 𝑣 𝑝 and particle diameter 𝑑 𝑝 :

𝑑 𝑚 = 4𝑎𝑑 𝑝 3𝑣 𝑝 (4)
where 𝑎 is a constant. Increasing the particle volume fraction and decreasing the particle size result in reduced grain size, thus leading to enhanced yield strength of metal matrix nanocomposites.

Coefficient of thermal expansion (CTE) and elastic modulus (EM) mismatching.

After solidification of the composite, the temperature keeps decreasing to room temperature.

Due to the difference of the coefficients of CTEs between the particles and the matrix, stresses are generated at the interface, leading to geometrically necessary dislocations (GND). GND can also generate from the difference of elastic modulus between particles and matrix during straining. The GND density due to CTE and elastic modulus can be calculated from the following equations: [START_REF] Sanaty-Zadeh | Comparison between current models for the strength of particulate-reinforced metal matrix nanocomposites with emphasis on consideration of Hall-Petch effect[END_REF] 𝜌 𝐶𝑇𝐸 = 𝐴?𝑎?𝑇𝑣 𝑝 𝑏𝑑 𝑝 (1-𝑣 𝑝 )

(5)

𝜌 𝐸𝑀 = 6𝑣 𝑝 𝑝𝑑 𝑝 3 𝜀 (6)
where 𝐴 is a geometric constant, ∆𝑎 and ∆𝑇 are the differences of the CTE and of temperature, respectively, 𝑒 is the bulk strain of the composite.

These models could be used to estimate the effect of different parameters of the composites. However, some critical factors are needed to make these models valid: a full density, good adhesion between the reinforcement particles and the matrix, and a uniform distribution of the particles. This is where the understanding and control of the process appear critical.

Plenty of materials have been commercially used as reinforcements. Among them, hard materials have often been chosen as reinforcement phase: covalent hard materials/diamondlike structure (C, CNx, BN, SiC, AlN); metallic hard materials/metal like structure (TiC, TiN, AlB2, WC, Ti2N, TiAlCN…) and ionocovalent hard materials/ionic structure (Al2O3, ZrO2…). [START_REF] Zhang | Recent advances of superhard nanocomposite coatings: A review[END_REF] Table 1.1 Commercially available nanoparticles used for reinforcement 37

Strategies to improve mechanical properties

Several strategies have been used to improve the mechanical properties through the mechanisms detailed above, such as selecting suitable reinforcement materials and optimizing particle characteristics (such as size distribution, composition and interfacial bonding), which will be discussed in the following paragraph.

Selection of suitable reinforcements

Recent studies on MMNCs focus on the light metal matrix. Among metals, the light alloys of Al, Ti and Mg are desirable as metal matrix. Their properties depend on the reinforcement particle features, more precisely, their morphology, size, load fraction, as well as the quality of the interface with the matrix. [START_REF] Li | Enhanced Mechanical Properties of Graphene (Reduced Graphene Oxide)/Aluminum Composites with a Bioinspired Nanolaminated Structure[END_REF] Carbon nanotubes (CNTs) are among the most studied strengthening additives for MMC. [START_REF] Bakshi | Carbon nanotube reinforced metal matrix composites -a review[END_REF] Their strong sp 2 bonds lead to high elastic stretching properties. For multi-wall carbon nanotubes, the theoretical elastic modulus can reach 600 and 1100 GPa while the tensile strength ranges from 35 to 110 GPa. [START_REF] Peng | Measurements of near-ultimate strength for multiwalled carbon nanotubes and irradiation-induced crosslinking improvements[END_REF] Bakshi and Agarwal 41 have reviewed the strengthening properties of CNT-Al composites. The record yield strength and elastic modulus are 621 MPa and 110 GPa, respectively. [START_REF] Choi | Reinforcing effects of carbon nanotubes in structural aluminum matrix nanocomposites[END_REF] Other allotropes like nanodiamond and graphene have also been studied for their loadtransfer ability and high thermal conductivity. Li et al. [START_REF] Li | Enhanced Mechanical Properties of Graphene (Reduced Graphene Oxide)/Aluminum Composites with a Bioinspired Nanolaminated Structure[END_REF] and good chemical inertness (oxidation up to 950 °C). They strongly improve mechanical properties and thermal conductivity. [START_REF] Lahiri | Insight into reactions and interface between boron nitride nanotube and aluminum[END_REF] These values varied with the processing technique.

BN nanoparticles and microparticles have also been embedded in an Al matrix. With powder spark plasma sintering (SPS) and ball-milling, Firestein et al. [START_REF] Firestein | High-strength aluminum-based composites reinforced with BN, AlB2 and AlN particles fabricated via reactive spark plasma sintering of Al-BN powder mixtures[END_REF] Carbides like TiC 47 , SiC, [START_REF] Gopalakannan | Application of response surface method on machining of Al-SiC nano-composites[END_REF][START_REF] Sujan | Physio-mechanical Properties of Aluminium Metal Matrix Composites Reinforced with Al2O3 and SiC[END_REF] TaC, WC, B4C [START_REF] Miracle | Metal matrix composites -From science to technological significance[END_REF] have been investigated by many researchers.

SiC is one of the most common reinforcements for MMNCs because of its high strength, low thermal expansion, high thermal conductivity, high hardness, high elastic modulus, high chemical inertness, and its quite good wettability by the Al matrix. Li et al. [START_REF] Li | Enhanced combination of strength and ductility in ultrafine-grained aluminum composites reinforced with high content intragranular nanoparticles[END_REF] Among all the ceramic particles used as reinforcements, titanium diboride (TiB2) has attracted the interest of many researchers because of its excellent resistance to mechanical erosion, higher hardness (960 HV), high melting point (2790 °C) and good thermal stability. Besides, TiB2 has been used as a grain refiner in Al alloys to suppress the growth of Al grains. [START_REF] Nampoothiri | A Review of Recent Progress in Solid State Fabrication of Composites and Functionally Graded Systems Via Friction Stir Processing[END_REF] Nanoparticles of TiB2 about 31 nm diameter were used to reinforce Al by cryomilling and spark plasma sintering (SPS). The strength of the composite was improved by 20% compared with equivalent SPS-consolidated 5083 Al alloy. [START_REF] Nampoothiri | A Review of Recent Progress in Solid State Fabrication of Composites and Functionally Graded Systems Via Friction Stir Processing[END_REF] Nampoothiri et al. 53 re-melted a TiB2/Al composite prepared by powder metallurgy and treated it with ultrasound process. The sizes of the TiB2 particles decreased after this treatment, yielding a two times increased yield strength compared to samples before treatment.

AlB2 54 and ZrB2 55 were synthesized by in situ reaction with boron and with a Al-K2ZrF6-KBF4 system. For example, Tian et al. 55 introduced pre-dried salts (K2ZrF6 and KBF4)

wrapped in an aluminum foil into molten 2024 Al alloy to promote the in situ reaction in the molten salt between the different precursors. The formed ZrB2 nanoparticles of ca. 30 to 100 nm were rectangular and were also uniformly distributed in the final composite. ZrB2reinforced nanocomposites exhibited higher tensile strength and yield stress.

f) Intermetallic materials

The review of Ward-Close and Doorbar 56 summarized the research on intermetallic matrices.

Intermetallic compounds like NiAl and Al3Ti 2457 have been successfully used as reinforcement phases in MMNCs. The Al-Al3Ti nanocomposite revealed good mechanical behavior at high temperature, and TiAl-NiAl MMNCs showed low fracture toughness and high hardness.

g) Combination of several reinforcement particles A combination of several reinforcement particles could produce a synergetic effect. For example, Sujan et al. [START_REF] Sujan | Physio-mechanical Properties of Aluminium Metal Matrix Composites Reinforced with Al2O3 and SiC[END_REF] studied the performance of stir casted Al2O3 and SiC reinforced metal matrix composite material. The result showed that the composite materials exhibit improved physical and mechanical properties, such as low coefficient of thermal expansion as low as 4.6 10 -6 •N - Non-uniform distribution of reinforcement objects is one of the main obstacles for enhanced mechanical properties. 58 Several techniques have been explored to circumvent this difficulty.

Ultrasound-assisted casting is effective to reach an homogeneous dispersion of the reinforcement objects in the matrix phase. 14,59,60 Another method consists in enhancing the The size of the reinforcement particles plays a critical role in the mechanical properties of MMCs. For example, a A356 alloy reinforced with nanoscaled TiB2 possessed higher elongation and tensile strength than with microscaled particles. 63 Another example from Youssef El-Kady et al. 64 deals with alumina-reinforced A356 alloy. Rheocasting and squeeze casting were used to incorporate 60 and 200 nm Al2O3 nanoparticles into A356 alloy. The nanocomposite with 60 nm Al2O3 particles showed higher ultimate tensile strength, yield strength and elongation compared to the composite with 200 nm Al2O3, while microhardness showed the opposite trend, probably because of the agglomeration of nanoparticles.

c) Volume fraction of the reinforcement phase

The content of reinforcement phase in the nanocomposite is also an important parameter to control the mechanical properties. and the ultimate compressive strength from 313 to 392 MPa, respectively. A365/Al2O3 metal matrix nanocomposites also showed improved ultimate tensile strength, yield strength and elongation with particle size of 60 nm. However, the excessive incorporation of nanoparticles could be detrimental to mechanical properties. 64 Indeed, Akbari et al. 63 found that increasing the content of TiB2 nanoparticles into A356 nanocomposites from 0.5 to 1.5 wt% significantly enhanced the final yield stress and ultimate yield strength, while an opposite trend appeared when the nanoparticle content kept increasing to 3.0 wt%. This behavior could be ascribed to increasing particle aggregation.

d) Interfacial bonding

A strong interface between the metal matrix and reinforcement objects is necessary to ensure effective load transfer. 68 Wang et al. 69 utilized reduced graphene oxide (RGO) to reinforce a Cu alloy (in presence of Ti) and found that a TiC layer was formed at the interface of RGO/CuTi composite. The TiC layer led to enhanced interfacial bonding enabling yield strength and tensile strength of 242 MPa and 307 MPa, 24 and 16% higher than those of 2 vol% RGO/Cu composite, respectively. Another example relates to the interface reaction between BN and Al that leads to the formation of AlN and AlB2. [START_REF] Lahiri | Insight into reactions and interface between boron nitride nanotube and aluminum[END_REF] Their formation in Al-based composites is not as deleterious as compared with carbides (Al4C3). 70 AlN reinforcement improves the hardness and flexural strength of Al matrix. AlB2 is also less detrimental to the fracture toughness of Al. [START_REF] Lahiri | Insight into reactions and interface between boron nitride nanotube and aluminum[END_REF] As for the previously described parameters, a tradeoff value exists for interfacial bonding. Indeed, excessive interfacial bonding may deteriorate the mechanical properties of the composite. Zhang et al. 21 applied spark plasma sintering to fabricate Ti/nanodiamond nanocomposites and found the formation of the new phase TiC derived from the reaction between Ti and nanodiamond at the grain boundaries. Small amounts of in situ formed TiC was regarded as one of the main strengthening mechanisms since it is likely to induce dislocation pinning. However, larger amounts of TiC (≥ 0.5 wt% nanodiamonds) tended to decrease the mechanical properties of the final nanocomposite.

Phase change materials for thermal management of Li-ion battery

As mentioned above, MMNCs are potential light-weight materials for thermal control, which is essential for vehicle electrification based on lithium batteries technology. Below, we are presenting major strategies for thermal control of batteries together with a brief overview of phase change materials. Electrical vehicles (EVs) including hybrid electric vehicles, plug-in hybrid electric vehicles and battery electric vehicles focus lots of attention because of their ability to reduce greenhouse gas emissions. 72,73 One of the main challenges to develop EVs lies in the energy storage system.

Importance of thermal management of the Li-ion battery

Lithium-ion batteries possess some advantages compared to other types of batteries, such as nickel-cadmium and nickel-metal hybrid batteries: a) higher energy density (Figure 1.4), b) higher specific power, c) lighter weight, d) lower self-discharge rates, e) higher recyclability and f) longer life cycle. 71,74 A full battery pack consists in several connected electrochemical cells (in series and/or in parallel). 71 A traditional cell contains an anode (graphite), a cathode (LiCoO2 layer), a separator and an electrolyte (Figure 1.5). On charging, the cobaltite is oxidized, leaving Li + , which goes through the separator and inserts into the interlayer space of graphite, which then is reduced.

The overall process implies an electric current and an electron transfer through the external circuit. 75 This reaction is reversed during discharge. However, the working temperature affects greatly the kinetics of these reactions, thus the performances, life and safety of Li-ion batteries.

Furthermore, temperature changes of Li-ion battery due to environmental conditions and heat release from several cycles of charging and discharging are inevitable. 73 In Tesla's Model 3 car, the battery pack could release more energy (in form of heat) than 6000 iPhone handsets. 76 Therefore, battery thermal management is necessary. We focus below on management of overheating. Heat generation within Li-ion battery is derived from resistive dissipation, relaxation of the concentration gradient in the cell and chemical reaction. 27,73 It can be simply described

Cooling strategies

The major role of battery thermal management is to maintain a suitable temperature range and evenly distributed temperature within battery during charging and discharging, even at extreme environmental conditions. 78 The most effective working temperature of Li-ion battery ranges from 15 to 40 ℃. 79 That means that the battery should on one hand reduce its temperature when operating temperature increases or environmental temperature is high, on the other hand increase temperature when the environmental temperature is low. The battery thermal management system itself should be compact, light, cost-saving, reliable and easy to operate, to pack and to maintain for its application in electronic vehicles. 80,81 Battery thermal management system could be categorized according to a) medium type, b) power consumption and c) whether it directly contacts with battery or not. 73 Here, we are going to discuss mediumbased cooling systems: air cooling, liquid cooling and phase change materials.

Air cooling. The simplest way to cool a battery is to blow air on the battery walls through a fan or a wind tunnel. In this case, during design of batteries, several parameters need to be taken into account, such as the battery layout, airflow or wind speed to minimize the energy consumption and costs. For example, Xu and He 82 improved heat dissipation by simply changing the longitudinal battery pack to a horizontal battery pack, which shortened the airflow path. Furthermore, they found that a double "U"-type duct enhanced heat dissipation performance rather than double "1" type. Around the battery, the air flow takes time to blow the whole system, leading to unevenly distributed thermal gradient. To solve this issue, bidirectional or reciprocating airflow has been implemented, which relies on cell temperature monitoring to actively adjust the airflow direction.

Liquid cooling. Although air cooling could achieve battery thermal management as described above, its efficiency is not sufficient to dissipate heat from an electric vehicle battery pack under a high charge/discharge rate at high working temperature, mostly because of the low thermal conductivity of air. for prismatic power batteries. As a control, pure paraffin enabled a dramatic reduction of battery surface temperature and maintained a suitable range of working temperature because of its latent heat absorption and of the natural convection of melted paraffin during melting event. A copper-paraffin composite could further reduce the surface temperature and enhance the uniformity of the distribution of temperature due to the improvement of thermal conductivity.

Metal matrix nanocomposite with phase change materials for temperature

Commonly used phase change materialsorganics and salt hydratesare restrained to applications at relatively low temperature and suffer from limited mechanical resistance, low stability upon cycling and low thermal conductivity. Researchers have incorporated thermally conductive fillers (such as graphite, metallic nanoparticles) and foams (metals and graphite, for instance) to produce phase change composites with improved thermal conductivity. The role of fillers is limited due to a discontinuous structure. Thermal interface resistance between phase change materials and fillers still prevail. On the opposite, a continuous structure with high thermal conductivity could be obtained after incorporating phase change materials into foams. 90 Due to the difficulties of infiltration process and to corrosion, foams are difficult to implement for phase change materials with high melting temperature. 91 Metallic materials have been found applicable as phase change materials for their high thermal conductivity (~10 to 400 W m -1 K -1 ), high melting temperature (-40 to 300 ℃) and large enthalpy of fusion. 92 Metallic phase change materials are then good candidates for thermal management.

To our knowledge, the first documented work on metal matrix nanocomposites with phase change inclusions was published in the 1990s. Goswami et al. 93 have incorporated Bi into glassy Al-Fe-Si matrix by mixing Bi with the matrix through induction melt spinning.

Nanocomposites containing Bi nanoparticles between 5 to 50 nm were achieved. This work was revisited in 2015, with the aim of improving the control of the nanoparticle size. Liu et al. 92 embedded Bi nanoparticles into an Ag matrix by directly adding the desired amount of Bi nanoparticles to molten Ag. The Ag matrix was found efficient to prevent coalescence of Bi nanoparticles and the phase change properties were stable even after 100 melt-freeze cycles.

Compared to common organic phase change materials, the incorporation of Bi nanoparticles resulted in 50-100% improvement of thermal energy density. In addition, by tuning nanoparticle size from 8 to 15 nm, the melting temperature of the silver matrix was shifted from 236 to 252 ℃. To convert into a less expensive aluminium metallic matrix, Abraham et al. 26 ball-milled a mixture of Al and Bi2O3. Subsequently, this pellet was heated and held over 300 °C to promote aluminothermic reduction of Bi2O3 and form an Al2O3 layer surrounding Bi particles within an Al matrix. The matrix encapsulating Bi inclusions could store thermal energy upon heat-freeze cycles and retained its structure. In the control sample with nonencapsulated Bi inclusions, molten Bi tended to separate from the Al matrix and to accumulate on the surface of materials.

MMCs fabrication techniques

Most manufacturing techniques for the synthesis of bulk metal matrix nanocomposites have already been reviewed in the literature among which we can cite the review of Bakshi et al. [START_REF] Bakshi | Carbon nanotube reinforced metal matrix composites -a review[END_REF] as well as the latest review of Casati and Vedani 24 . Metal matrix composites/nanocomposites can be produced by many techniques (Table 1.2) mainly composed of solid-state, liquid-state and semisolid-state processes. 2 In the following paragraphs, friction stirring, spark plasma sintering, ultrasound-assisted casting, vortex casting, squeeze casting bottom-up fabrication and electrodeposition are going to be discussed.

Table 1.2 Class and process techniques of powder metallurgy 10

Powder metallurgy

Powder metallurgy includes all the production processes using metal powders as matrix precursor. It consists of three main steps: mixing metal matrix powders with the reinforcement phase in the form of powders, consolidating the powders mixture, post-treating the mixture to trigger e.g. sintering. [94][95][96] The aluminum and/or reinforcement powders are obtained by ball milling. Mechanical alloying, a repeated process of cold welding, fracturing and re-welding of powder particles in a high-energy ball mill, is a fundamental technique for the preparation of MMNCs. 11 Several techniques have been used to perform the compaction of composite powders: hot pressing, cold pressing, cold isostatic pressing (CIP), hot isostatic pressing (HIP), equal channel angular pressing (ECAP) followed by a sintering treatment.

Powder metallurgy produces well-dispersed metal matrix composites with uniform microstructure. It can be used to produce high melting point metals like Ti with less difficulty than liquid metallurgy. However, many treatment steps are necessary, which increases the price of fabrication. The presence of clusters of nanoparticles is also difficult to prevent. The products also show a limited range of shapes that is detrimental to vehicle industry. Therefore, the focus is put in this PhD work on liquid processing. Nonetheless, we describe below some powder metallurgy processes that are the main approaches used up to now to reach MMNCs.

Friction stir processing is one of the powder metallurgy approaches. The introduction of reinforcement objects into aluminum matrix can also be performed by multi-pass friction stir process. The principle is based on three major steps: firstly, fixing the matrix sample;

secondly, drilling holes according to a predesigned pattern with pins into the matrix phase;

thirdly, inserting nanoparticles into the holes. 97 Multiple passes are necessary for a uniform distribution of reinforcement particles. The friction stir process enables the introduction of more than one reinforcement. For example, Jalilvand et al. 97

Liquid state process

In liquid metallurgy techniques, the reinforcement phase is added to the molten metal matrix in order to attain a desirable composition. There are some significant challenges to incorporate nanoparticles in molten metals. Firstly, the melt temperature is relatively high, some nanoparticles are not stable enough in these conditions; second, nanoparticles tend to float on the surface of the molten metal due to the poor wettability of the nanoparticles with the molten metal; third, nanoparticles agglomerate due to attractive Van der Waals force and to the lack of a repulsive force in molten metals; and fourth, nanoparticles are normally pushed to the grain boundary by solidification fronts during solidification. 103 Three main liquid state processes are discussed here: stir casting, vortex casting, squeeze casting. Stir casting has been widely applied as a commercial technique to fabricate MMCs. In conventional stir casting, the metal matrix is firstly de-moistured and heated above its melting temperature. Then the de-moistured reinforcement phase is added into the molten metal, followed by stirring with a stirrer to homogenize the reinforcement/metal matrix mixture. After cooling down in a mold, the final MMCs with desired shape could be obtained. But two main challenges prevail for this technique: one is related to the quantitative incorporation of the objects in the melt, the other one is the homogeneous distribution of the nano-reinforcements because of the high surface-to-volume ratio of the objects and of their poor wettability by the molten matrix, resulting in particle segregation and floating due to the density difference between the matrix and the nano-reinforcement objects. 104,105 In order to solve these challenges, ultrasonication has been developed to disperse nanoparticles in molten metals (Figure 1.8). Eskin 106 has explained the application of ultrasound treatment to light alloy melts. In melts, the mechanism of cavitation is similar to other liquids. High intensity ultrasonic waves are introduced to generate transient cavitation and acoustic streaming, which are used to refine the microstructure, degas the melt and finally homogenize the distribution of reinforcement nano-objects. 107 During this process, the ultrasonic cavitation produces a strong implosive impact that could on one hand help to well disperse the nano-reinforcements and on the other hand enhance the wettability of the reinforcements by the molten metal. 108 The varied material-dependent propagation sound speed (see Table 1.3), thermal conductivity and viscosity influence the acoustic impedance and the attenuation factor. The surface tension of metal melts is significantly higher than that of water (0.86 N m -1 for Al at 700 °C vs. 0.079 N m -1 for water at 20 °C), which decreases the dispersion efficiency. In order to achieve a good dispersion and distribution of nano-reinforcements, this other difficulty must be overcome when fabricating MMNCs. 106 The production of composite materials with particles via ultrasound-assisted casting comprises several steps.

Ultrasound-assisted casting

Step 1: Preparation of the particles with suitable surface treatment to decrease the surface tension and increase wetting of the reinforcement phase particles by the molten metal.

A coating with metal like Cu or Ni, oxidation (like Al2O3), degassing and drying is useful to reach an homogeneous dispersion. 24 Step 2: Preparation of the metal matrix. Additives for surface tension control are useful, especially for the liquid Al melt, which has a high surface tension. A lower surface tension metal like Bi, Pb, Sb, Li, Ca, Mg and Sn can be used to decrease the surface tension of Al. 109

Step 3: Incorporation of the particles. The simplest technique consists in spraying the particles onto the surface of the melt. This technique works relatively well for magnesium

MMCs 19 but has limitations in for MMCs due to the thick oxide film at the surface of the melt and to the high surface tension of the melt. In some cases, mechanical mixing using an impeller under protective atmosphere are combined with ultrasonic mixing. 110 Besides, the ultrasound processing parameters, the melt temperature, duration of processing, size, morphology and content of reinforcement particles can be adjusted to improve the dispersion. 24 One way to incorporate small particles while preventing flotation and macroscopic phase separation is to add the particles to the matrix in the semiliquid state (the temperature of the metal matrix is put between its liquidus and solidus temperature). For instance, SiC nanoparticles were added to a slurry of Mg-18% Zn alloy and dispersed with mechanical stirring. Then the composite material was heated to the liquid state and ultrasonicated to get a satisfying dispersion. 111 Step 4: Casting and solidification. Ultrasonication in light alloy melts may affect nucleation and release of dissolved gas from the liquid phase. In semiliquid mixture, it influences the structure formation by fragmentation of solid crystals and their distribution.

In order to well understand the effect of ultrasonic stirring on the structure and mechanical properties of MMNCs, Jia et al. 112 applied ultrasonic treatment to incorporate 1 wt.% SiC or Al2O3 nanoparticles to aluminum alloy A356. The authors firstly melted the alloy in an induction furnace and then added nanoparticles into molten liquid under ultrasonic stirring with power of 1.75 kW, frequency of 18 kHz. Subsequent solidification was performed to obtain aluminum metal matrix nanocomposite. The ultrasonic treatment could change the grain structure from dendritic to globular shape and also contribute to reasonable dispersion of nanoparticles into the final composite although some agglomeration still existed.

Furthermore, the tensile strength of the nanoparticle-reinforced composites was higher with ultrasonic treatment.

The ultrasound assisted casting is highly dependent on the stability of nanoparticles and wettability of the inclusions by the matrix. Several combinations 1 are proven to be efficient, especially for the SiC/Mg6Zn couple. Research of novel nanoparticles/metal matrix is still in need to achieve better mechanical properties or other atypic properties.

Vortex casting

Vortex casting has been proven useful to uniformly distribute reinforcement objects into a metal matrix. In a typical experiment, the milled reinforcement phase and matrix powders are continuously added to molten aluminum or aluminum alloy under stirring. Another stirring is performed to homogenize the mixtures after several minutes' crystallization. The final MMCs are obtained by shaping in a mold and an extrusion process. However, this method is difficult to be used for the incorporation of nanosized particles as it is difficult to reach high dispersion.

Squeeze casting

Squeeze casting is also a casting technique that involves two main steps: a) reinforcement objects addition into the molten metal matrix under mechanical stirring; b) solidification under high squeeze pressure. This technique can reduce porosity and shrinkage, improve casting yield and enhance properties of the final MMCs. 113 Lü et al. 60 firstly fabricated a millimeter-sized granule composed of 8 wt% SiC nanoparticles in aluminum by dry high-energy ball milling, then remelted and diluted this granule in molten A356 alloy to get final 1 wt% nanoparticles composites under ultrasonication. The molten mixture was poured into a die cavity that was disposed into an environment with 200 MPa squeeze pressure for 30 s until the end of solidification (Figure 1.9). As a result, the nanoparticles were uniformly distributed into the nanocomposite and tensile strength of squeeze casted MMNCs was 2169 MPa, 25% higher than the non-reinforced alloy matrix. then sintered to decrease the porosity. This bottom-up approach allows a perfect dispersion of the nanocrystalline phase, which could increase hardness and strength. However, porosity still exists and the fabrication needs pre-synthesized nanoparticles without passivating oxide layer, which is difficult to achieve for highly reductive metals like magnesium and aluminum.

Electrodeposition

The group of Akbulut presented a new hybrid nanocomposite thermal interface material (TIMs) relying on copper or silver, boron nitride nanosheets (BNNSs) and organic ligands. The composite was prepared by chemisorption-coupled electrodeposition approach. 115 BN nanosheets were exfoliated and functionalized by bifunctional ligands, which were supposed to link the metal and the BNNSs together through Lewis acid-base interactions. The coelectrodeposition/chemisorption process was then applied, which coupled the electrolytic reduction of metal ions and the chemisorption of functionalized BN nanosheets on the cathode.

The thermal conductivity of the final composite was 236-258 W m -1 K -1 and 306-321 W m - 1 K -1 for physically and chemically conjugated BNNS-silver-ligands materials, respectively, while their hardness varied from 400-495 MPa and from 240 to 360 MPa, respectively. The electrodeposition method is suitable to fabricate coating materials, but is hard to be applied to bulk materials.

Conclusion and motivation of this work

Faced with energy and environmental challenges and in line with the consumer demand, the Stellantis has been investing for several years in manufacturing vehicles that are more envir onmental friendly. To execute this strategy, the group wishes to invent and fabricate new materials that are light-weight and low cost with high ductility and also with good thermal conductivity. In this context, this PhD work aims to investigate MMNCs with improved mechanical properties and thermal properties. In particular, the fabrication processes must be designed to meet the needs of the current casting process of alloy.

Among the four types of fabrication techniques shown in this chapter, ultrasonicassisted casting can meet these requirements. The use of ultrasonic treatment is a good way to improve the dispersion of nanoparticles during the casting process of aluminum alloy. This process also offers advantages such as low cost, easy molding process and adaptability to largescale fabrication.

The literature review of current MMNCs opens many intriguing opportunities for the development of new materials. Regarding the mechanical properties of MMNCs, commercial inorganic nanoparticles have been proven useful to efficiently improve strength and functionality of the metal matrix. The reinforcement nanoparticles must be designed to satisfy the strengthening mechanisms and harsh casting conditions. Boride nanoparticles synthesized by molten salts are good candidates as they are small in size, light, hard and stable at high temperature. At the time when we started this project, there was only two papers available in the literature that used molten salt synthesized nanoparticles (particles size between 10 to 100 nm) as ex situ reinforcement. 63 Additionally, the working temperature is also limited: phase change only occurs at the melting temperature of pure metal, which is too high in most cases. To overcome these issues, we propose a liquid phase method: metallic nanoparticles (pure metal or alloy) will be encapsulated in an inorganic shell to prevent macroscopic phase separation in the melt.

Subsequently, these nanoparticles will be incorporated in molten metal matrix by ultrasoundassisted process. This work will be presented in two parts: chapter 2 will deal with the synthesis of the building block core-shell nanoparticles, chapter 4 will deal with the fabrication of the nanocomposites and their thermal properties.

Introduction

As described in chapter 1, metal matrix phase change nanocomposites are promising candidates for the thermal control of battery packs. To design such nanocomposites, we have selected Bi nanoparticles as proof-of-concept phase change material for their relatively low melting temperature. We focus on an aluminium matrix to target light weight composites. A Bi@Al2O3-Al nanocomposite has been successfully prepared previously by powder metallurgy. 1 In their process, a mixture of Al and Bi2O3 was ball milled then the pellet was heated and held above 300 °C to promote aluminothermic reduction to form an Al2O3 layer surrounding Bi particles within the Al matrix. The Main limitation of its use on battery pack is its high fabrication cost and low adaptability of the temperature range. Therefore, we investigate herein the ability of liquid metallurgy processes to design metal matrix phase change nanocomposites. To do so, we must first elaborate an approach that will avoid the coalescence of molten Bi droplets in the molten aluminium matrix. For this purpose, we seek to protect the Bi nanoparticles against coalescence by a silica shell to prevent nanoparticles coalescence. The objective of this chapter is to present the synthesis and characterization of these building blocks: bismuth nanoparticles encapsulated in a silica shell.

State of the art on the Bi synthesis

Bismuth nanoparticles are of interest for several applications such as catalysis 2 and biomedical contrast agents in computed tomography scan or X-Ray imaging. 3 In the latter case, they are especially good candidates due to their low toxicity and high biocompatibility. 4 The synthesis of bismuth nanoparticles has been performed by using the sol-gel process, 4 In order to expand the range of diameter size, the authors modified the synthesis protocol by tuning concentration ratios between Bi[N(SiMe3)2]3, PHD-co-PVP and NaN(SiMe3)2, leading to a tuning of the particle size over a wide range between 3-115 nm and with various morphologies such as Bi nanorods and nanoplates. 4 The authors also applied a hot injection method to reach a diameter ranging from 30 to 45 nm. But the synthesis procedure still shows some limitations: 1) the necessity of an additional step to prepare the Bi[N(SiMe3)2]3 precursor, and 2) the unavailability of the stabilizing polymer PHD-co-PVP. 7 BiCl3 can also be used as Bi precursor to obtain Bi nanoparticles with size range of 4-29 nm. These particles were proved as good catalysts for solution-liquid-solid growth of high-quality CdSe quantum wires. 5 A simpler method has been developed by Scheele et al., 8 by reducing Bi(OAc)3 with oleylamine (OA) or trioctylphosphine (TOP) in the presence of thiols as stabilizing ligands at lower temperature (60 °C). The resulting Bi nanoparticles were obtained as monodisperse samples, especially for a diameter of 3.3 ± 0.5 nm. The size of the nanoparticles could be increased from 7-9 nm to 10-14 nm by tuning the ratio between thiol and oleylamine, and could be further increased by replacing oleylamine with TOP to get nanoparticles of 40 nm.

solvothermal
A stable suspension of Bi nanoparticles without toxic organic residue is one of the main challenges for the application of Bi nanoparticles in the field of biomedicine. Solvothermal synthesis and hydrothermal synthesis have been developed afterwards. Wang et al., 9 applied solvothermal synthesis approach for Bi nanoparticles with different shapes. NaBiO3 was reduced by ethylene glycol in presence of the capping polymer of poly(vinyl pyrrolidone) (PVP) or a trace amount of Fe 3+ at 200 ℃. By tuning the ratio of PVP and Bi, also the amount of Fe 3+ , the authors could achieve different shapes for the crystals such as spheres when the PVP/Bi ratio was 0, and nanocube when the ratio was 1.6. In presence of a trace amount of Fe 3+ , the PVP/Bi ratio of 1.6 resulted in nanobelted crystals. Microwave heating could also assist this approach.

Theoretically, hydrothermal synthesis is difficult to be applied to fabricate oxidatively prone elements such as bismuth, which is prone to get oxidized in water. Brown et al. 10 developed an aqueous synthesis of Bi nanoparticles based on a pH-dependent solubility equilibrium. An aqueous solution of Bi(NO3)3•5H2O was reduced by NaBH4 at pH 7.4 in the presence of dextran as a stabilizer. The nanoparticles were stable between pH 7 and pH 12.

Increasing the synthesis pH inside this range could slow down the reaction rate and lead to more dispersed Bi nanocrystals. However, the shape of nanoparticles was not well controlled.

The main reported methods for the synthesis of Bi nanoparticles have been presented

above. In our work, we were particularly interested in obtaining bismuth nanoparticles smaller than 40 nm, since these particles exhibit a depression of the melting point around 250 °C, lower than the melting point of bulk Bi at 271 °C, 11 which could enable a decrease of the phase change temperature in the final metal matrix nanocomposites. The quantity of the product in each batch is also a critical factor because these nanoparticles will serve as precursors for the synthesis of nanocomposites, which requires more than 100 mg of nanoparticle for each lab-scale (see chapter 3) synthesis. Solvothermal and hydrothermal methods are attractive in this respect, but the samples are highly polydisperse and these syntheses often result in large particle size over 40 nm. We then selected the method developed by Scheele et al. 8 as it could satisfy the requirements described above.

State of the art on the growth of silica shell to encapsulate bismuth particles

The synthesis of silica coated metallic nanoparticles has been well studied for nanoparticles such as silver and gold. 32 Among the numerous applications of metal@silica coreshell nanoparticles, most biological applications of nanoparticles require a good control of the surface state of nanoparticles. The implementation of a silica shell helps to graft specific reactive functions such as antibodies or fluorophores. 32 Among all synthesis methods, we should cite two main processes: a) the Stöber process, a sol-gel process where the Si precursor (typically tetraethyl orthosilicate (TEOS)) is hydrolyzed and condensed in an alcoholammonia-water solution; 33 b) the use of a reverse microemulsion before condensation. [START_REF] Hull | Introduction to dislocations[END_REF] The growth of a silica shell on Bi nanoparticles has been less studied in the literature.

Lee et al. [START_REF] Sanaty-Zadeh | Comparison between current models for the strength of particulate-reinforced metal matrix nanocomposites with emphasis on consideration of Hall-Petch effect[END_REF] developed Bi nanotube@SiO2 core-shell structures by atomic layer deposition. Chen et al. [START_REF] Zhang | Recent advances of superhard nanocomposite coatings: A review[END_REF] reported Bi@SiO2 nanospheres of 100 nm via in situ reduction of Bi2O3@SiO2 by NaBH4. Luo et al. [START_REF] Malaki | Advanced Metal Matrix Nanocomposites[END_REF] synthesized around 80 nm Bi@SiO2 with 5 to 20 aggregates encapsulated in the same silica shell. The synthesis method was not described in the article. Li et al. [START_REF] Li | Enhanced Mechanical Properties of Graphene (Reduced Graphene Oxide)/Aluminum Composites with a Bioinspired Nanolaminated Structure[END_REF] synthesized 40 -60 nm Bi nanoparticles in acidic solution at 80 °C, then the authors grew a 2 -5 nm SiO2 layer on the Bi nanoparticles by the Stöber process. Li et al. [START_REF] Bakshi | Carbon nanotube reinforced metal matrix composites -a review[END_REF] succeeded to achieve monodisperse core-shell objects. Spherical Bi nanoparticles (70-75 nm) were synthesized by thermal reduction of Bi(NO3)3 in dodecanethiol at 178 °C. Then, a 40 nm-thick SiO2 shell was grown by employing the reverse-microemulsion method: Bi nanoparticles were dispersed in cyclohexane and treated with ultrasonication, then polyoxyethylene, nonylphenylether and ammonia solution were added and stirred for 24 h.

In view of the facility of the synthesis and of the quantities of product achievable, the Stöber process has been applied in our work.

Synthesis and characterization of Bi nanoparticles

Firstly, we applied the protocol developed by Scheele et al. 8 for the synthesis of Bi nanoparticles. This protocol consists in reducing bismuth acetate in oleylamine in the presence of 1-dodecanethiol as ligand at 60 °C following a hot injection protocol, the experimental details are described herein: The powder X-ray diffraction (XRD) patterns confirmed that Bi nanoparticles were obtained (Figure 2.1a). TEM images (Figure 2.1b) showed that the sample consists in bismuth nanoparticles with diameter around 25 nm with a relatively narrow size distribution. Contrary to the literature in which nanoparticles of 12-14 nm were obtained by this process, 8 the nanoparticles we synthesized are larger and the sample is more polydisperse. We then evaluated the impact of a range of parameters to understand this difference. In brief, the particle size distribution did not exhibit significant dependence on the origin and purity of the reagents, on the synthesis temperature (60 ± 2 °C), neither on the reaction duration from 20 to 24 and 28 h.

When the reaction time was decreased to 6 h, the nanoparticles (figure 2.1c) were more faceted.

The only remaining synthesis parameter is the injection speed and/or the heating ramp after the injection. These parameters have not been described in the original article of Scheele et al. 8 Then, we have attempted to preheat oleylamine at 55 °C before injection. The resulting nanoparticles shown in figure 2.1d exhibit smaller particle size ranging from 4 to 24 nm with a wide range of morphologies compared with the sample prepared by hot injection of roomtemperature oleylamine (figure 2.1b). Because the final nanoparticles are dedicated to the design of phase change materials relying on the melting temperature decrease of nano-sized bismuth, nanoparticles smaller than 30 nm are sufficient for this study as this size range corresponds to a noticeable decrease of the melting temperature. [START_REF] Peng | Measurements of near-ultimate strength for multiwalled carbon nanotubes and irradiation-induced crosslinking improvements[END_REF] Thus, we selected nanoparticles of 25 nm diameter (figure 2.1b) from room temperature oleylamine injection for the following steps.

Synthesis and characterization of Bi@SiO 2 nanoparticles

In a second step, the bismuth nanoparticles washed with 2 centrifugation-dispersion cycles with ethanol and chloroform and were coated by a thick silica shell via the Stöber method. [START_REF] Cao | Study on tensile properties and microstructure of cast AZ91D/AlN nanocomposites[END_REF] The protocol is as follows:

Tetraethyl orthosilicate (TEOS) was dissolved in absolute ethanol at a concentration of 0.38 mol L -1 . 10 μL of this diluted TEOS solution and 100 μL of ammonia (30%) were subsequently added to 1 mL ethanolic bismuth suspension at [Bi] = 6.10 -4 mol.L -1 under vigorous stirring. The mixture was heated to 40 °C for 3 hours. After cooling down, it was washed two times with ethanol.

57

The silica shell obtained is clearly observed on TEM images, with a lower contrast than that for the Bi core (Figure 2.2a). Its thickness was around 7 nm. TEM (Figure 2.2a) shows that the silica shell actually embeds several particles, forming aggregates with size up to 160 nm.

Since the shape of aggregates is irregular, here we characterize its size with the area distribution (Figure 2.2c). The area varies from ~1500 nm 2 to ~20000 nm 2 . The area of a single core nanoparticle is ~ 1500 nm 2 , and the area of the largest object observed correspond to the 14 times the size of single object. Since nanoparticles partially overlap in TEM image, the largest objects possess ca. 20 nanoparticles inside the same silica shell. This aggregation phenomenon is usual for the encapsulation by the Stöber process of poorly stable ethanolic suspensions of small nanoparticles. [START_REF] Hull | Introduction to dislocations[END_REF] The number of washing cycles of the initial particles strongly influences the size of the final objects as well. With subsequent centrifugation and washing steps, nanoparticles are more aggregated and yield Bi@silica aggregates larger than 1 micrometre after 5 washing steps (figure 2.2b). One explanation could be the removal of the stabilizing ligands on the surface of nanoparticles after washing. This phenomenon has been well described in the literature for gold nanoparticles, [START_REF] Choi | Reinforcing effects of carbon nanotubes in structural aluminum matrix nanocomposites[END_REF] in which repeated washing removed the ligands, yielding irreversible aggregation. As repeated washing results in aggregation of nanoparticles, no more than two washing steps were applied in order to preserve the suspension's stability in ethanol. We then attribute the formation of Bi@SiO2 aggregates during silica growth to the low dispersibility of Bi nanoparticles in ethanol. By taking inspiration from the case of gold nanoparticles, [START_REF] Ma | Nanometric Si3N 4 particulate-reinforced aluminum composite[END_REF][START_REF] Chaira | Synthesis of Aluminium -Cementite Metal Matrix Composite by Mechanical Alloying Synthesis of Aluminium -Cementite Metal Matrix Composite[END_REF]32 we applied ligand exchange to incorporate a functionalized thiol molecule that would ensure higher dispersibility of the particles in ethanol in the ligand monolayer. The ligand exchange protocol is described below:

After the synthesis of Bi nanoparticles and two washing cycles with ethanol and chloroform, the particles were redispersed in ethanol. A given quantity of ligand, varying from 0 to 5 ligand per Bi surface unit (nm 2 ), was solubilized in ethanol and was added into the reaction medium, followed by heating to 40 °C for 1 h in a closed vial in an oven.

At the end of their synthesis, the pristine Bi nanoparticles were covered by 1dodecanethiol as ligand. Mercaptoundecanoic acid (MUA) and mercaptohexadecanoic acid (MHA) (Table 2.2) were assessed to perform ligand exchange. The thiol group is expected to graft on the surface of the nanoparticles and the acid group is expected to improve the dispersion ability in ethanol via hydrogen bonding. Different lengths of the spacer chain between MUA and MHA were used to tune the hydrophobicity of the ligand and of the nanoparticles.

MUA MHA

The quantity of ligands is also inspired from the case of gold nanoparticles. The total coverage ratio of thiol on gold surface was 5 thiol•nm -2 . [START_REF] Chaira | Synthesis of Aluminium -Cementite Metal Matrix Composite by Mechanical Alloying Synthesis of Aluminium -Cementite Metal Matrix Composite[END_REF] Series of experiences have been performed to study the influence of the exchanged ligand on the aggregation state of Bi cores within the silica shell/matrix. In order to verify the efficiency of ligand exchange before growth of the silica shell, we observed the bismuth nanoparticles via TEM just after ligand exchange (in ethanol) and then after redispersion in an apolar solvent (hexane). Here, we should notice that the TEM grids were prepared by drop casting of nanoparticles' suspensions onto carbon-coated copper grids, which triggers further aggregation during evaporation. Nonetheless, we could not detect significant differences in terms of dispersibility between bare particles and ligand exchanged ones (figure 2.4a-c): aggregates of around 15 nanoparticles down to single nanoparticles were both observed in each sample. The effect of ligand exchange was not significant in ethanolic suspensions. After redispersion in hexane, we observed aggregates around 200 nm, 400 nm and even larger than 800 nm for the samples without ligand exchange, exchanged with MUA and exchanged with MHA, respectively (figure 2.4d-f). The increased aggregation state with exchange of ligands indicated that ligand exchange actually occurred, but the increase of the hydrophobicity of the ligand didn't improve the dispersion in hexane, which could be ascribed as the influence of organic residuals as observed in TEM images. The aggregation state observed from the pristine ligand-exchanged Bi nanoparticles was maintained after growth of the silica shell (figure 2.5). The difficulty to reach a stable suspension of Bi nanoparticles is then the bottleneck that hinders isolation of single Bi@SiO2 nanoparticles. This could arise from the reaction medium's composition (e.g. side product residuals from the Bi nanoparticles synthesis) and from intermolecular interactions between thiols, thiolates and Bi-thiolate complexes, which probably show differences in complexing abilities on different metal surface.

Conclusion

In this work, we succeeded in synthesizing core-shell Bi@SiO2 nanoparticles with 25 nm Bi core and 7 nm silica shell. Clusters of several particles are embedded in the silica, forming aggregates of size up to 160 nm. By tuning the number of washing cycles of the pristine Bi nanoparticles, the most dispersed samples contain 1 to 15 nanoparticles that are encapsulated in the same silica shell. Single Bi nanoparticles encapsulated in silica shells have already been achieved for particles larger than 40 nm by the reverse-microemulsion method [START_REF] Bakshi | Carbon nanotube reinforced metal matrix composites -a review[END_REF] and by acidic solution synthesis followed by Stöber method. [START_REF] Li | Enhanced Mechanical Properties of Graphene (Reduced Graphene Oxide)/Aluminum Composites with a Bioinspired Nanolaminated Structure[END_REF] The encapsulation of Bi particles smaller than 30 nm by silica is still a challenge. Compared to previous results on silica-encapsulated Bi nanoparticles with particle size below 30 nm, which showed aggregates of 5 to 25 nanoparticles [START_REF] Malaki | Advanced Metal Matrix Nanocomposites[END_REF] and more than 20 nanoparticles 36 encapsulated in the same silica shell, our protocol shows enhanced dispersion of the Bi cores, although we could not obtain only single cores. Nevertheless, these nanoparticles meet the requirements described in the introduction: the particle size is small enough to trigger a decrease in the melting temperature and the synthesis was successfully scaled-up to 80 mg of product per batch. Another unexpected advantage is that the large size of the silica aggregates could facilitate the identification of the reaction between silica and aluminium. This part will be discussed in detail in chapter 4.

Introduction

One of the key challenges in fabricating metal matrix nanocomposites by liquid metallurgy processes lies in the harsh conditions of the nanoparticles feeding and dispersion process at high temperature (over solidus or liquidus of aluminum alloys, generally over 550 °C). Indeed, the stability of nanoparticles at high temperature should always be kept in mind in order to obtain the final material. In the past decades, classic aluminum matrix nanocomposites have relied on commercial nanoparticles with a limited choice. 1 Recent progress in the in situ growth of oxide, boride and carbide nanoparticles has shown some promises in the increase of nanoparticle loading and in the interfacial bonding between nanoparticles and metal matrix. 2,3,4 In particular, the reaction of K2TiF6 amd KBF4 in molten mixtures of aluminum and fluoride salts enabled the in situ formation of TiB2 nanoparticles in different aluminum alloys. 4 However, two main concerns still exist: one is related to the difficulty to control the particle size and the other one lies in the presence of impurities in the final products. 4,5,6,7,8 The continuous development of nanoparticle synthesis approaches, especially the emergence of molten salts synthesis offers an alternative solution. 9 Several classes of nanoparticles including borides, nitrides, carbides and silicides have been synthesized using molten salt solvents. 10 The use of inorganic molten salts as high temperature solvents (300 to 900 °C) could lead to nanoparticles that are stable at the processing temperature of liquid metallurgy. As-synthesized nanoparticles possess wellcontrolled particle size and they could be easily dispersed in water or alcohols after several washing. All these developments taken together, nanoparticles prepared in molten salts appear as interesting potential precursors to create new metal matrix nanocomposites. Nevertheless, there are still several issues to overcome to reach this aim, such as the demixation of nanoparticles powders from the molten matrix during their incorporation and during solidification.

In this work, we aimed at designing boride-aluminum nanocomposites in the presence of 99.9% pure aluminum and of an aluminum alloy AlSi7MgCu (Table 3.1) by synthesizing ex situ the nanoparticles. Boride nanoparticles are hard and light-weight and are promising candidates to further reinforce light-weight alloys. In this chapter, we attempt to establish the suitability of boride nanoparticles synthesized in molten salts as a potential reinforcement for aluminum-based matrices. HfB2 was selected as a target because it is a superhard material, with a predicted Vickers hardness of 30.5 GPa, 11 It also contains hafnium, which is expected to facilitate the electronic microscopy observations and chemical detection in the final composites and will then be a useful probe to reach a proof-of-concept material. The dispersion behaviour of the nanoparticles in molten aluminum and its effect on the enhancement of mechanical properties are evaluated. The results are then compared with those obtained for reference SiCbased nanocomposites, a well studied aluminum-based matrix system.

Noticeably, at the time when we started this project, there were only two articles available in the literature that used molten salt synthesized nanoparticles (particles size between 10 to 100 nm) as external reinforcement. 12 But other related works have been published recently, especially from Li Group 13 who worked on the incorporation of molten-salt synthesized TiB2 and ZrB2 nanoparticles in aluminum. However, none of them have applied hafnium boride nanoparticles as reinforcement so far.

Alloying elements of alloys for cast parts

Materials and methods

Preparation of nanoparticles

HfB2 nanoparticles were prepared in molten salts as previously reported. 15 0.5 mmol of hafnium (IV) chloride, 1 mmol of sodium borohydride and 2.5 g of eutectic mixture LiCl/KCl powders were ground together with a Retsch MM400 ball miller for 2 minutes at 20 Hz. The reaction medium was heated in a carbon crucible under argon flow at 900 °C for 4 h. After cooling down to room temperature, the nanoparticles were washed with 12 cycles of centrifugationredispersion in water.

The synthesis of amorphous boron nanoparticles was similar to HfB2 nanoparticles. 6 mmol of sodium borohydride was grounded with 5 g of eutectic LiI/KI. The medium was heated at 800 °C for 1 h. SiC nanoparticles with a diameter of 35 nm were purchased from Nanomaker, France.

Preparation of nanocomposites

HfB2-Al and SiC-Al nanocomposites were prepared by ultrasound-assisted casting of a mixture composed of nanoparticles and aluminum. Two sets of experiments were designed. The first approach involved a direct introduction of nanoparticles into the molten metal. The metal matrix was heated at the desired temperature for 5 minutes. The desired amount of nanoparticles powder was packed in an aluminum foil and dipped into the molten metal at 720 °C. A steel spatula was used to stir the melt for 2 minutes. Then, a high-power ultrasonic probe made of niobium with tip diameter of 4 mm (QSonica sonicator Q500A-220, USA) was applied to disperse the nanoparticles for 5 minutes, with an interval of 5 s on / 5 s off at 30% of the nominal power. Finally, the sample was cast in an alumina crucible. For characterization, the middle part of the sample was cut and analyzed.

In the second approach, 2 g of aluminum powders (325 mesh, 99.99%) were added into an suspension containing the desired quantity of nanoparticles powders into 2 mL of ethanol, dispersed by sonication at 80 °C as setpoint for 1 hour. Then the mixture was dried in a vacuum oven at 40 °C for 1 h, followed by pressing into a pellet under 8 tons for 2 minutes.

Subsequently, the pellet and 1 g of pure Al beads were heated by an induction oven at 720 °C under argon atmosphere in an alumina crucible. After dwelling for 1 min, ultrasounds were applied as described above to disperse nanoparticles for 5 minutes, with an interval of 5 s on / 5 s off. The final composites were obtained by cooling down the melt in the crucible. They were then cut by diamond wire saw for further analyses.

X-ray diffraction

X-ray diffraction was performed with a D8 Bruker diffractometer operating with CuKɑ radiation in Bragg-Brentano configuration. The ICSD reference card 30422 was used for the indexation of the HfB2 structure.

Transmission electron microscopy

For transmission electron microscopy (TEM), the nanoparticles were dispersed in ethanol, then one drop was evaporated on carbon-coated copper grids. For the prepared composites, TEM thin foils were prepared by a rectangle of 3*3*0.5 mm cut in the middle of the sample. The polishing step was then performed by diamond disc to get a 200 μm thick sample and the thickness of the centre was further reduced to 50 μm by using a precision dimpling instrument (Model 515, South Bay Technology). Ion milling was performed for the final thinning. TEM images were recorded on a Tecnai spirit G2 microscope operating at 120kV. The STEM images and STEM-EDX mapping were recorded with a JEM 2100Plus UHR microscope, and were both performed at 200 kV. The size distribution of nanoparticles was determined by counting 150 to 200 nanoparticles from TEM images.

Scanning electron microscopy

The microstructure of the nanocomposites was characterized by scanning electron microscopy (SEM). The samples were mounted in polyfast resin (Struers), followed by a series of polishing steps through 1000, 2400, 4800 grit papers (Buehler) then with 3 µm and 1 µm diamond suspensions. For SEM observations, the samples were directly attached on a SEM holder using carbon tape. They were then observed under scanning electron microscopy (SEM, Hitachi S-3400N and Hitachi SU-70 microscope) equipped with an Oxford X-Max 50 mm2 EDS detector.

Image analysis was performed with the Fiji software. The grain size distribution was measured by ImageJ from optical microscope images and SEM images.

Dynamic Light Scattering

The average hydrodynamic diameter of HfB2 aggregates dispersed in water was determined by dynamic light scattering (DLS) in the Zetasizer Nano ZS90 (Malvern Panalytical).

Microhardness tests and nanoindentation tests

The Vickers hardness was measured with a load of 0.05 kg, dwell time 10 s using Wilson VH1102 microhardness testers.

The nanoindentation measurements were realized using a Nano Indenter XP® equipped with a Berkovich tip. The hardness was measured at the unloading process using Oliver and Pharr method. 400 measurements were achieved on a matrix of 21*19 with a distance of 10 µm at a depth of 1 µm to obtain a mapping of mechanical properties. The distance between two scratches was 10 µm. A matrix of 6*5 with a depth of 3 µm was performed to obtain an average measurement of mechanical properties.

Results and discussion

Preparation and characterizations of nanoparticles

HfB2 nanoparticles were prepared by molten salts synthesis. 15 Different contrasts observed by TEM (figure 3.1a) confirm the core-shell structure of the nanoparticles. TEM shows a diameter of the contrasted core of 4 to 14 nm. This core shows high contrast, which is ascribed to a crystalline character and to a higher electron density, due to the Hf content. The powder XRD pattern (figure 3.1b) is fully indexed along the hexagonal HfB2 structure. The crystallite size of HfB2 calculated according to the Scherrer formula is 8 nm, in agreement with TEM, thus indicating that the core is single crystalline and made of HfB2. A less contrasted layer of ca.

3 nm shown covers the surface of HfB2 nanoparticles, corresponding to a layer of amorphous boron as described in a previous study. 15 The diameter measured by DLS on an aqueous dispersion is about 400 nm, which is much larger than the TEM data, indicating the formation of aggregates. These objects are clearly observed by TEM, which indicates that the amorphous shell embeds several particles. According to X-ray photoelectron spectroscopy (XPS) published in a previous work, 15 the boron surface is partially oxidized for the nanoparticles washed with water. that SiC is the only crystalline phase.

Preparation of metal matrix nanocomposites

To overcome the difficulties of directly loading the nanoparticles powders in the melt (see Appendix S3.1), which are mainly ascribed to the poor wettability and high volume of nanoparticles in powder form, we used a two-step process to introduce nanoparticles in the melt: a) dispersing the nanoparticles in ethanol in the presence of 325 mesh aluminum micropowder, then pressing the slurry into a pellet; b) diluting the pellet in molten aluminum under ultrasonication.

Pressing of the pellet at room temperature without heating prevented the coalescence of the nanoparticles and excluded the potential reactivity between Al micrograins and the boron shell of the nanoparticles before the dispersion process. 12 The experimental setup is shown in figure 3.2a. Ultrasonic waves generate cavitation and acoustic streaming, which is most likely to refine the microstructure, break aggregates and homogenize the distribution of nanoparticles. 16 However, most commercial ultrasound microprobes for < 5mL volume are made by Ti alloy. This material is not applicable in our case because of the erosion of the Ti rod in molten aluminum 17 and the reactivity with the nanoparticles, especially their boron shell. A niobium tip of 4 mm diameter has been specially

designed by the supplier company to limit the contamination of the samples (figure 3.2b). We further designed the shape of the crucible, as shown in figure 3.2c, to increase the dipping depth of the ultrasonic tip. The tip was dipped in 2 mL of molten aluminum at a depth of 3 mm, corresponding to 60% of the depth of liquid Al. This set-up is suitable for preliminary studies without the need to scale up the synthesis of nanoparticles. Furthermore, the ultrasound treatment power and intervals have also been tuned.

Previous works have shown that the mechanical properties of composites prepared by discontinuous ultrasonic treatments were better enhanced than with the continuous ultrasonic treatments due to enhanced dispersion. 16 In our study, an interval of 5s on/ 5s off was applied for all experiments. Considering the liquid splash at high power and pore formation at intermediate power, we chose to work at an intermediate ultrasound power at 30% of the maximum power followed by 30 s degassing with a low power of 10%. These conditions were then set as a model for the following experiments.

More details for the improvement of the process can be found in appendix S3.1.

Microscopy analysis of HfB 2 and SiC nanocomposites HfB 2 -Al nanocomposites

Several experiments were performed by dispersing 0.8 vol.% HfB2-Al nanoparticles into aluminum using conditions defined in 3.3.2. The experimental parameters are presented in table 3.2 (P83). Broad peaks from XRD patterns (Appendix 3.5) confirmed that the phase was maintained. Low intensity peaks at 28° and 41° might be indexed along the AlB12 phase, which may arise from the reaction between molten aluminum and the boron-rich shell of the nanoparticles.

These samples were further analyzed by electron microscopy coupled to EDS. TEM images (figure 3.4) also showed that nanoparticles formed into sub-micron aggregates at the Al grain boundaries and no separated nanoparticles could be found in Al grain.

Compared to the initial nanoparticle precursors, HfB2 nanoparticles could maintain their size after the harsh melting process at high temperature (figure 3.4a) although their morphology slightly changed from sphere-like particles to faceted particles (figure 3.4b). The HfB2 crystal structure was confirmed according to the lattice spacing observed by HRTEM, also indicating the thermal stability of HfB2 nanoparticles. The mismatch between other HfB2 facets and Al lattice planes, except for the plane couple described above, could disturb crystalline order of Al near the interface, and thus cause Al lattice distortions near the interface.

The combination of possible epitaxial relationship (a) and lattice distortion (b) implies that the HfB2-Al interface could be semi-coherent, which is expected to favour the mechanical properties. 18 The modification of the shape of the nanoparticles implies also that the particles undergo facet reconstruction, favouriting (010) facets epitaxially oriented versus (100)Al planes. c) Amorphous-like defects. (yellow arrow, within yellow dashes) This could be attributed to the reactivity of the matrix with the ca. 2 nm thick amorphous boron-rich layer exhibited by the initial nanoparticles. The size range of this area at the interface between HfB2

and Al in figure 3.4b is consistent with the thickness of the pristine layer. According to the Al-B phase diagram (Appendix 3.6), 19 AlB2 and AlB12 could form above 500°C. Although the kinetics of the reaction between nanoscaled boron (in the form of the nanoparticles' shell) and molten Al are unknown, a previous study showed that AlB2 AlB10 and AlN form at the interface between BN nanotubes and aluminum after 10 minutes at 650°C, 20 which is corresponding to conditions similar to those used in our experiment. The Gibbs free energy for the formation of aluminum boride from Al and boron is similar to the one for the reaction described in the literature, 20 so the presence of crystalline aluminum boride is highly possible in our sample. No fringe is observed, probably because this phase is not well oriented.

The identification of the nature of the boron-rich shell is necessary to differentiate boron from the Al matrix. It is a strong challenge as boron cannot show contrast differences in the presence of Al. We performed high resolution TEM in a region where the Al matrix was not well oriented: Al fringes cannot be observed (Appendix 3.7). A crystalline layer around 2 nm

shows lower contrast compared with the 4 nm dark core. The HRTEM micrograph cannot be indexed satisfactorily along structures such as Al, Al2O3, Hf, HfO2, HfB2, HfB, AlB2, AlB10, but could be attributed to AlB12, in agreement with XRD (Appendix 3.5).

SiC-Al nanocomposites

A 0.8 vol% SiC-Al nanocomposite was obtained via the approach detailed in the previous section. It showed results similar to HfB2-Al nanocomposites: nanoparticles mostly segregated at the grain boundaries according to SEM images (figure 3.5a-e). Figure 3.5f shows a TEM image of a SiC nanoparticle in a 200 nm Al grain. This indicates that part of the SiC particles are located inside the grains, witch is according to the previous literature, 16 and contrary to our TEM observations for HfB2-Al composites.

In fact, one of the main challenges for the preparation of metal matrix nanocomposites is that nanoparticles tend to segregate towards the grain boundaries at the solidification fronts upon cooling down. This phenomenon can be explained by several reasons: a) in both SiC and HfB2 samples, nanoparticles form aggregates in the solid state due to attractive Van der Waals forces, while the applied ultrasound cavitation and acoustic streaming in molten metal is not sufficient to disperse aggregates into single particles. HfB2 are also initially located into a boron-rich layer that embeds irreversibly several particles together. b) during solidification, attractive Van der Waals forces and interfacial forces overcome the Brownian motion and viscous drag force, 21 leading to particle aggregates, c) SiC and Al have a rather large equilibrium lattice mismatch, 22 which is detrimental to interfacial stability and energetically disfavors an homogeneous dispersion, d) The presence of a passivation oxide layer onto Al, as

shown by EDS mapping, may also impact the wetting behaviour of Al on SiC.

Previous works have been performed to improve the dispersion of nanoparticles within

SiC-Al alloy nanocomposites. 23,24,25 A better dispersion has been achieved by Lu et al., where diluted granules are prepared by powder metallurgy and dispersed into a molten Al alloy by ultrasonication. Although micro-sized clusters were observed, nanoparticles were not totally pushed to the grain boundaries. 21 An even better dispersed nanocomposite could be achieved by changing the composition of the matrix to tune the wettability and the rheology, such as for the dispersion of SiC in Mg-Zn alloys. 26 The optical microscopy images (figure 3.6a) highlight the Si phase in dark grey colour.

SEM-EDS mapping (figure 3.6b-g) suggests that hafnium is located in the grain boundaries rich in Si phase. According to the precedent results, it is not surprising that nanoparticles were segregated to the grain boundaries.

We also investigated the effects of the nanoparticles on the structure of the metal matrix during solidification. A pure aluminum pellet was added to AlSi7Mg0.3Cu0.5 alloy and prepared by the same method (Appendix 3.9). Comparing the sample without treatment (Appendix 3.10) and the sample treated by ultrasounds (Appendix 3.9), we observed globular grain structures after ultrasound treatment. This phenomenon is in accordance with previous reports. 1,16,27 When the nanoparticles were incorporated to the alloy, the microstructure of the alloy was slightly refined, with Al grains of an average size of 90 μm. The globular shape was also modified.

However, this behavior is less marked than with HfB2-Al samples. As a matter of fact, the solute present in the alloy is responsible of constitutional undercooling during solidification, which plays a primary role in decreasing the size of the grains. The grain size in the alloy is then dominated by the composition of the alloy. As described earlier, wettability and viscous drag force are the key factors that may maintain the nanoparticles dispersed in the metal grains and enable them to resist the solidification front. 28 However, the control of the freezing conditions has also an influence on the microstructure. Increasing the cooling rate can form a refined network of nanoparticles by quenching the samples, thus leading to high modulus for the metal matrix grains. 29 Therefore, the microstructure was assessed for 0.8 vol% nanoparticles-Al nanocomposites obtained by three different cooling methods: a) For samples heated in the resistive furnace, metallic samples are always inside an alumina crucible during cooling. Relatively slow cooling is achieved by air cooling of the quartz tube. The solidification time is 2 minutes from 720°C to 660°C. The cooling from operating temperature to room temperature takes 2 hours (around 6°C min -1 ).

b) For samples heated in the induction furnace, only the metallic melts are heated, the crucible is only heated by contact with the Al melt. As the heating time is short (10 minutes), we expect that the crucible is colder than in the case of the resistive furnace and the cooling rate is then higher. The molten metal returned to the solid state in around 10 s after heating. The cooling from operating temperature to room temperature takes approximately 40 minutes (around 18°C min -1 ).

c) Another series of samples was obtained by resistive heating and cooling relatively quickly by immersion in a liquid nitrogen bath. The sample was cooled to the temperature of liquid nitrogen (no more bubble appears) in 10 minutes, with a cooling rate of around 100 °C min -1 .

Firstly, the homogeneity of the microstructure was investigated. In conventional air freezing, the microstructure of the sample on the top, middle and bottom section was homogeneous. Nucleation occurs homogeneously inside the sample and Al grains grew in an anisotropic fashion. For the samples cooled by liquid nitrogen, presented in Appendix 3.11, the microstructure at the core of the sample remains homogeneous, but the top and the bottom areas

show completely different morphologies. The homogeneous core might be due to the presence of an important temperature gradient between the core and the surface of the sample. It is accepted that liquid on the surface solidifies almost immediately when the sample is plunged in liquid nitrogen. Top view images showed large grains around 60 μm. In this region, the liquid metal was in contact with the alumina crucible, which is a low heat conductor. Fewer nucleation sites are formed at the low cooling rate region. The nucleated grains grow into a macroscopic grain during solidification. Thus, the 100 μm layer on the surface and in contact with the crucible are not considered for further analysis.

Analysis of the average grain size reveals a refinement of the microstructure with the increase of the heating rate. Indeed, 0.8 vol% HfB2-Al samples cooled at different rates showed an average grain size of 120, 78 and 20 µm for cooling rates 6, 18 and 100°C min -1 , respectively.

(figure 3.7) SiC-Al nanocomposites also follow the same trend. Liquid nitrogen cooling is therefore proven to be the most efficient among the three methods assessed to refine the microstructure of MMnCs. A common cooling method in the industry is water quenching. The difference between water quenching and liquid N2 quenching has been studied in the course of this PhD work (appendix 3.12). Regarding the average grain size for different cooling methods, liquid nitrogen quenching showed similar grain refinement behavior (82 µm with natural cooling, 48 µm for water cooling and 41 µm for liquid nitrogen cooling). Considering the wide distribution of grain size, the difference between these two cooling methods was not significant. Even though the temperature of liquid nitrogen is much lower than that of water, the formation of N2 bubbles at the surface of hot metal inhibits heat transfer and then decrease the cooling ability of liquid nitrogen, thus explaining the similar microstructure of water and N2-cooled samples. Water cooling and liquid nitrogen cooling could both be applied with variable sets of parameters.

Influence of the inclusions load on the microstructure of nanocomposites

Aiming at improving the dispersion of nanoparticles, we compared a series of composites prepared at identical conditions (5 minutes ultrasound treatment by Nb tip, cooling by liquid nitrogen) but with different inclusion loads.

Optical microscopy images (Figure 3.8) indicated a surprising increase of grain size with the inclusion load. In line with previous SiC-Al nanocomposites without extra dispersion methods, 16 a volume ratio of 1.5 % is hard to achieve. This phenomenon might be explained by 82 macro-segregation observed e.g. in the samples at 2.4 vol% of inclusions. This phenomenon could relate to the viscosity and its impact on the dynamics of the solidification front during the solidification process. Therefore, a ratio of 0.8 vol% was kept as a model.

Mechanical tests

Microhardness measurements

The mechanical properties were firstly investigated by microhardness measurements and then compared to the pure reference samples prepared in the same way (table 3.2). For all experiments, the hardness at grain boundaries (rich in nanoparticles) was much higher than the value within the grain, as expected from the analysis of microstructure. Considering the inhomogeneity, an average of measurements on grain boundaries and inside the grains was also calculated. For nanocomposites containing 0.8 vol% of nanoparticles (SiC and HfB2) sonicated for 30 min and quenched by liquid nitrogen, the measured hardness (table 3.2) for nanocomposites was notably greater than the hardness of the base material: average HvSiC = 40 HV and HvHfB2 = 33 HV compared to 24 HV in pure Al. The increase of the hardness of > 40% is likely due to the presence of hard nanoparticles at the grain boundaries, which enhances the hardness of the final composites. In order to understand in more depth the mechanisms underlying hardness increase, it is worth characterizing average hardness values inside Al grain. The average Hv inside the Al grains for the sample 0.5vol% SiC-Al is 32.2 HV, higher than the value for pure Al (24 HV), this notable increase is more likely to be attributed to the presence of SiC modulus of nanocomposites at the sub-microscopic level. The results were compared with reference samples prepared in the same way in the absence of nanoparticles. As described in the previous section, SEM images and microindentation showed that the ultrasound treatment time has no significant effect on the microstructure and the hardness of the final material for both SiC and HfB2 nanocomposites, and the hardness is largely related to the location of aggregates of nanoparticles. This allows us to compare samples prepared with different ultrasound treatment times. The mapping of the mechanical properties correlated with the optical image of the samples is helpful to distinguish the interior of the grains and grain boundaries. 400 measurements were achieved on a matrix of 21*19 with a distance of 10 µm at a depth of 1 µm (figure 3.9). The hardness of several areas could not be reliably measured (black pixels in the hardness maps and values equal to 0) due to the existence of micropores, especially for the sample at 2.4 vol% SiC-Al. These defects could be ascribed to the scratching of nanoparticles during polishing steps, or the presence of micropores in the sample. In fact, with the increase of the load, Al melt tends to be more viscous and micropores are more difficult to be eliminated, further investigations on SEM should be performed to confirm this. Figure 3.9 presents the mapping of mechanical properties of the 0.8% HfB2-Al nanocomposite. A heterogeneous hardness distribution was observed: a population with a relatively homogeneous hardness corresponding to the grains and a second population significantly higher for the grain boundaries. Similar observations were made for other nanocomposites: SiC-Al and HfB2-AlSi7Cu0.5Mg0.3 (Appendix 3.14). The high hardness/modulus regions (grain boundaries) and low hardness/modulus regions (grains) will be discussed separately in the following sections. 3.3). The 0.8 vol% SiC-Al composite exhibited higher hardness than the 0.8 vol% HfB2-Al sample, with values 0.52 and 0.49 GPa, respectively (table 3. The grain matrix and the intergranular regions were then compared. Measurements at 1 µm depth revealed the reinforcement effect at the extreme surface. In low hardness regions, the middle of the plateau for 0.8 vol% SiC was higher than that of pure Al (0.6 GPa vs 0.43 GPa) (figures 3.10b and 3.11b). Hardness increases in the low hardness regions, correspond the matrix grain boundaries, implies trapping of part of the SiC nanoparticles inside Al grains, which is in accordance with TEM observations described above. Conversely, the hardness distribution for the 0.8 vol% HfB2-Al at 1 µm depth (figure 3.10b) shows low-hardness regions with hardness values similar to pure Al around 0.45 GPa, confirming that most HfB2 nanoparticles were segregated towards grain boundaries. The distribution curve for HfB2-Al (figure 3.10b) deviates towards higher hardness than pure aluminum for more than 50% of measurements (from 200 to 400), which corresponds to reinforced area at grain boundaries according to mapping (figure 3.9). The average of this population is 0.66 GPa (calculated from the highest 200 measurements), clearly higher than pure Al. These measurements indicated that the reinforcement effect of HfB2 nanoparticles occurs only at grain boundaries. 3.3), significantly higher than for the 0.8 vol% SiC-Al sample. The hardness of 2.4 vol% SiC nanocomposites has been increased by 68% compared to the reference aluminum matrix. However, we should also notice that with the increase of inclusion content, the standard deviation is also higher. The extremely high hardness regions up to 3 GPa were found to be aggregates of SiC. This indicates that the microstructure is more heterogeneous, as observed in microscopy images (figure 3.8a).

Effect of nanoparticles on the hardness of foundry Al alloy. In the case of the AlSi7Mg0.3Cu0.5 foundry alloy, a similar hardness has been observed: 0.79 GPa on average with 0.8 vol% HfB2 nanoparticles, compared with 0.77 GPa for the reference (table 3.3). The large standard deviation of 0.2 GPa was mainly related to the different hardness values of Al grains and of Si-rich regions in the alloy. The distribution curves (figure 3.12) were drawn from 30 measurements at 3 µm (figure 3.12a) and 400 measurements at 1 µm (figure 3.12n). The latter show an overlapping between the bare matrix and the HfB2-containing one in the 240 lowest hardness values. On the opposite, a deviation towards higher hardness values is observed for the highest 160 values, a similar fashion as for the Al matrix composites. The tangential midpoint of the distribution curves between the 250 th and 400 th measurements are 1.8 GPa for the sample with nanoparticles and 1.4 GPa for the reference matrix. Therefore, HfB2

Conclusion

Herein we have incorporated silicon carbide nanocrystals and HfB2 nanoparticles synthesized in molten salts in aluminum-based matrices by ultrasounds-assisted liquid metallurgy. Overall, the nanoparticles are stable after heating and aggregated at the grain boundaries. This segregation contributes to a refinement of the microstructure by decreasing the matrix grain size. Mechanical characterization by micro-and nano-indentation delineated the hardness and strength profiles of these composites. The hardness is heterogeneous at the sub-microscale, exhibiting a population with relatively homogeneous hardness corresponding to the matrix grains and a second population with significantly higher hardness for the intergranular zones.

The reinforcement effect is less pronounced for HfB2-based nanocomposites than for SiC-based samples, which could be ascribed to larger segregation of the boride particles at the grain boundaries. Indeed, inside the matrix grains, SiC nanocomposites are harder than pure aluminum, while the HfB2-based nanocomposites show mechanical properties similar to the Al reference. This further confirms the poor dispersion of HfB2 particles into aluminum, which is probably due to the initial aggregation state of the starting nanoparticles. Quenching the sample in liquid nitrogen allows refining the microstructure to limit the growth of the aluminum grains.

Other quenching methods, previously developed to fabricate amorphous alloys 30 should be explored in the future to obtain better dispersed HfB2-Al nanocomposite, as well as modified molten salt strategies to design less aggregated boride nanoparticles.

improved cooling efficiency and decreased local heat accumulation. 4,5,6 Abraham et al. 7 have further tried to solve the problem of weakened structural strength when the inclusions melt, by creating an Al matrix nanocomposites with Bi inclusions encapsulated with Al2O3. The composite with inclusions encapsulated in alumina shell presents higher yield strength compared to the non-encapsulated sample, and is capable of retaining its structure during cycling. Such encapsulation strategy would then be efficient to increase the mechanical resistance of phase change materials. However, this strategy relies on the reduction of metal oxide (Bi2O3) by aluminum, which shows severe limitations. First, the working temperature is limited, as phase change only occurs at the melting temperature of bulk low melting point metal like Bi, Ga, In, Sn. Second, the composite has been fabricated by powder metallurgy which requires long preparation time and high fabrication costs. Further investigations are then needed to develop an alternative encapsulation method and to move the fabrication process to liquid metallurgy in order to reduce the cost.

To alleviate those constraints, we propose to fabricate metal matrix nanocomposites with phase change inclusions by ultrasonic-assisted melt processing: casting the metal matrix over its melting temperature, then dispersing the encapsulated low melting-point metal in the liquid metallic matrix by ultrasonic-assisted treatment. 8 The dispersion of core-shell Bi@SiO2 in pure Al is studied herein as a proof of concept to assess the feasibility of the strategy and to study the reactions occurring between the nanoparticles core, shell and metal matrix under these harsh conditions. Here, aluminum was selected as a matrix due to its light weight, high thermal conductivity and high mechanical properties. Pure Al was used instead of industrial alloys in order to facilitate the understanding of the reactivity between the metal matrix and the inclusions. As Bi nanoparticles (melting point around 250 °C) are molten during the incorporation into the molten metal matrix (melting point 660 °C) and then show a high risk of coalescence, they were encapsulated in a silica shell prior to the composite preparation to prevent coalescence. The choice of silica as protective shell is motivated by polymer-based phase change materials, in which nanoencapsulation of phase change materials by silica has been proven to be an easy and efficient way to control the size of inclusion particles. 9 The synthesis of Bi@SiO2-Al nanocomposites consists in three steps: synthesis of nanoparticles, growth of a silica shell and incorporation in molten metal matrix by ultrasonic-assisted process, the samples were mounted in polyfast resin (Struers), followed by a series of polishing steps through 1000, 2400, 4800 grit papers (Buchler), then with 3 μm and 1 μm diamond suspensions.

For the observation of nanoparticles in transmission electron microscopy (TEM), the nanoparticles were dispersed in ethanol, then a drop was evaporated on a carbon-coated copper grid. For the composites, TEM thin films were prepared by a parallelepiped rectangle of

Chemical analysis by Wavelength dispersive X-ray fluorescence spectrometry (WDXRF)

Wavelength dispersive X-ray fluorescence (WDXRF) spectroscopy was performed on the cut pieces of the samples with a Bruker S8 Tiger spectrometer equipped with a Rh tube.

Differential Scanning Calorimetry (DSC)

The DSC analysis was performed in a TA Q20 DSC instrument. In a typical experiment, Around 10 mg of nanoparticles were sealed in an aluminum crucible. In the case of nanocomposites, top, middle and bottom slices of the samples (around 40 mg) were analysed separately to assess the homogeneity of the casted samples. To study the phase transition, the samples were heated to 320 °C and then cooled to 20 °C with a slope of 10 °C•min -1 followed by 3 minutes of stabilisation at 20 °C between each cycle. For each sample, At least 3 cycles were performed.

Results and discussion

Synthesis and characterization of Bi and Bi@SiO 2 nanoparticles

The first step of the composite preparation consists in the synthesis of size controlled bismuth nanoparticles (Figure 4.1), which was performed according the colloidal synthesis previously reported by Scheele et al. 10 The powder X-ray diffraction (XRD) patterns confirmed that Bi nanoparticles were obtained ( In a second step, the bismuth nanoparticles were coated with a 7 nm thick silica shell by the Stöber method, 11 as demonstrated by TEM, with a lower contrast for the shell than for the Bi core (Figure 4.2c). TEM (Figure 4.2c) shows that the silica shell actually embeds clusters of several particles, forming aggregates of size up to 160 nm. This aggregation is usual for the encapsulation of small nanoparticles by Stöber process, as suspensions of nanoparticles smaller than 100 nm are poorly stable and show a strong tendency to aggregate in an alcoholic solution. 12 To demonstrate the stabilizing action of the silica shell against coalescence of the bismuth cores, we have heated the core-shell particles and sole bismuth nanoparticles without inorganic shells above the melting point of bulk Bi (271 °C). The bare Bi particles are not stable when heated at 320 °C and undergo extensive coalescence and growth. (Appendix 4.2) In the opposite, the silica shell preserves the Bi cores from coalescence and ensures that the Bi core size distribution is maintained (Figure 4.2d). Although Bi nanoparticles are aggregated inside the silica shell, their size has been preserved when heated over the melting point, thus allowing us to use them in further steps to study the stability of Bi nanoparticles in molten aluminum.

Synthesis and characterization of Bi@SiO 2 -Al nanocomposite

Nanocomposites made of an aluminum matrix embedding Bi@SiO2 (Bi@SiO2-Al binding energy for Si-O bonds of SiO2. 16 We observe an additional contribution between 154 eV and 150 eV that is not observed for pure aluminum. This contribution is suggests the presence of other silicon species at lower binding energy than for SiO2, which is consistent with low oxidation state Si species and/or an aluminosilicate. 17 Note that oxidation of low oxidation state Si species could also occur during the preparation of the sample for XPS measurements and exposure to the air. Hence, we can hypothesize that silicon may be more reduced in the pristine sample than indicated by XPS.

This high-temperature phase may be induced by intense local heating provided by ultrasound treatment. 20 The formation of an aluminosilicate should be accompanied with the silicon, although no crystallized silicon was observed by TEM. The amorphous-like area we have highlighted (green square in STEM-EDS mapping in figure 4.4a,b and figure 4.5a) is then attributed to amorphous silicon. Note that this area could also correspond to crystalline silicon with an orientation away from diffraction conditions. Further analysis needs to be performed to confirm this.

Overall, we propose the following reaction mechanism: the pristine silica shell is reduced by aluminum that is concomitantly oxidized and reacts with remaining silica to form an aluminosilicate intermediate. During ultrasound-assisted heating for the nanocomposite preparation, the silica shell separates from the Bi core during, leading to leaching of Bi droplets which then undergo Ostwald ripening and develop into two populations.

Synthesis and characterization of Bi-Al nanocomposite

A comparative study of uncoated Bi nanoparticles dispersed in molten aluminum was performed. Figure 4.7 displays the HAADF and TEM images of such a Bi-Al composite.

Similar to the observations made for the Bi@SiO2-Al composite, the particle size evolved towards two populations due to Ostwald ripening: coalesced particles around several micrometres and detached particles around 5 to 30 nanometres. Thus, the particle size was no longer maintained at the nanoscale. The comparison between Bi-Al and Bi@SiO2-Al samples show that although the silica shell underwent chemical reaction and local mobilisation, it acted efficiently against coalescence of the cores occurring in the sample without protective shell. Consequently, a nanocomposite with wide size distribution at the nanoscale was generated.

Heat absorption ability evaluated by DSC measurements DSC analysis for nanoparticles

To assess the heat absorption ability, DSC was performed for both precursor nanoparticles and final nanocomposites. Figure 4.8 displays the melting behaviour of the 3 rd heating cycle for pristine Bi and Bi@SiO2 nanoparticles. Bi@SiO2 showed an endothermic peak at ∼242 °C, lower than the melting temperature of bulk Bi at 271.3°C, which is in agreement with the depression of the melting point with particle size decrease. 21 The melting point measured for 25 nm silica coated Bi nanoparticle is in accordance with the results of Bi-carbon nanocomposites. 22 Contrary to the similar melting behaviour for Bi@SiO2 nanoparticles upon cycling (figure 4.8b), the bare bismuth particles displayed a small exothermic peak at ∼245 °C, followed by two sharp endothermic peaks at ∼263 °C and 271 °C after 3 cycles, the latter corresponding to bulk Bi melting. The first two endothermic peaks are attributed to the melting temperature depression with the particle size. 21 The appearance of the peaks at higher temperature indicates that the nanoparticles coalesced and large particles up to bulk were We did not detect by DSC the phase change of small nanoparticles around 8 nm expected from TEM observations, which occurred at around 240 °C in Ag matrix. 21 This could be explained by the decreased enthalpy with the particle size 24 or the low volume fraction of nanoparticles smaller than 10 nm. These hypothesis need to be studied by further experiments.

DSC further confirms the efficient role of the silica shell to limit the coalescence of Bi nanoparticles, since the Bi-Al nanocomposite showed three endothermic peaks at 262, 269 and 271 °C overlapping together and shifted to higher temperature compared to the Bi@SiO2-Al nanocomposite, which implies the progressive coalescence of nanoparticles into bulk particles during cycling.

Compared to a Bi-Al composite prepared by powder metallurgy 7 at the same heating rate and with micron and nano-sized bismuth inclusions, the distribution of melting temperatures for Bi@SiO2-Al prepared by our liquid metallurgy process is broader and shifted to lower temperatures. This demonstrates a decreased Bi particle size achieved through our protocol. The small initial nanoparticle size, the protective silica shell, the ultrasound treatment induced disruption of Bi droplets and rapid thermal treatments by induction heating are the keys to achieve small inclusion nanoparticles with a wide phase transition temperature range, which is a considerable advantage for the final application as they can provide progressive heat absorption in a wide temperature range.

Conclusions

In conclusion, we have successfully developed an ultrasound assisted casting synthesis of phase change metal matrix nanocomposites. The phase change nanoparticles were embedded within a silica shell, which acts as a barrier against extensive coalescence of the metal cores, which are molten during the incorporation process.

We have made several attempts for in situ fabrication of Bi/Al2O3/Al by ultrasoundassisted casting method (see Appendix S4.1 and S4.2). The formation of encapsulated inclusions by aluminoreduction proved difficult to be applied to other types of inclusions such as In/Al2O3/Al: the reaction conditions must be adjusted with the Gibbs free energy for the formation of metal oxide. An additional challenge of prime importance relates to the decrease of the melting point of the PCM nanoparticles. Reaching other phase transition temperatures with metal alloy inclusions will require an indepth work in the future. We believe that such light-weight aluminum-based functional materials with high heat dissipation ability, wide and tunable phase transition temperature and low fabrication costs will be a relevant model from the point of view of thermal control of vehicle battery.

The electrical vehicle is expanding due to the gradual awareness and national policies of reducing fuel consumption and emissions of pollutants. But until now, there are still challenges that must be addressed to improve its performances and enable wide spread application. Long endurance as one important concern that requires the decrease of the vehicle's weight.

Electrification of the whole vehicle also encounters several challenges, particularly with regard to heat regulation. Improving the thermal management of batteries is a particularly important issue to ensure safety on board, to optimize cell aging and to preserve its performances. It is on these topics that the work in this thesis focused: fabrication of a metal matrix nanocomposites HfB2-Al with enhanced mechanical properties, and fabrication of a metal matrix nanocomposite with phase-transition inclusions made of phase change Bi cores encapsulated in silica, as a proof of concept for thermal regulation. The fabrication process applied in this work is liquid metallurgy assisted by ultrasound treatment that could be scaled up and implemented in current work flows for the production of aluminum alloys.

We have addressed the incorporation of nanoparticles synthesized in molten salts into aluminum-based matrices by ultrasound-assisted liquid metallurgy. An experimental set-up was designed and the process was carefully engineered to achieve total incorporation of nanoparticles in the matrix with decreased porosity and homogeneous chemical composition over the whole volume at the micron scale.

Several experiments were then performed by dispersing in molten aluminum 0.8 vol.% of HfB2 nanoparticles synthesized in molten salts. Commercial SiC nanocomposites were also assessed as references. According to electron microscopy, both types of nanoparticles are able to maintain their size and morphology after the high temperature melting process, but nanoparticles were segregated at the grain boundaries. HfB2 particles were only located at the boundaries, but SiC nanoparticles were also observed inside the Al grains, indicating better dispersion. Liquid nitrogen quenching was then applied to refine the microstructure. Parameters regarding the improvement of wettability (composition of the matrix, type of the nanoparticles, coating of the nanoparticles) or the quenching step (in particular other protocols developed to manufacture amorphous alloys) might be explored to gain further control over the dispersibility.

The strengthening effect of the nanoparticles was studied with nanoindentation tests.

The mapping of the mechanical properties correlated with the optical image, which showed heterogeneous hardness at the micro scale: a population with a relatively homogeneous hardness corresponding to the matrix grains and a second population with significantly higher hardness for the intergranular zones, in line with the high concentration of nanoparticles of hard phases at grain boundaries. The same phenomenon is also observed in other nanocomposites such as SiC-Al, HfB2-AlSi7Cu0.5Mg0.3. The average hardness of the sample with the inclusion of HfB2 nanoparticles was improved by 14%. This value is lower than that of SiC nanocomposite with the same volume ratio (21%), probably due to the difference in segregation state. The elastic modulus measured by nanoindentation showed no difference between pure Al and nanocomposites with 0.8 vol% loading of nanoparticles (SiC and HfB2). The tensile strength might also be measured to reveal the differences of elastic modulus at the macroscale.

Three main pathways were explored to fabricate bismuth-aluminum composites by liquid metallurgy. Firstly, Bismuth oxide was dispersed in molten Al and further reduced in situ into metallic bismuth by aluminoreduction. This process was proven hard to be replicated for other inclusions with are more difficult to be reduced, such as In or metallic alloys. Secondly, bismuth nanoparticles were directly incorporated in the Al matrix. Nanoparticles underwent large extent Ostwald ripening during the dispersion in molten Al, forming two groups of particles with diameters below 10 nm and about several micrometers, respectively. Thirdly, we have developed an encapsulation method by designing core-shell Bi@SiO2 nanoparticles, to incorporate in molten Al. The nano-inclusions also underwent size distribution widening towards smaller and larger sizes by Ostwald ripening, but to a smaller extent than the nonencapsulated sample since the maximal particle size of the enlarged particles was about only 200 nm. This result indicates that although the silica shell underwent chemical reaction and local mobilization, it acted efficiently against core coalescence. We have also demonstrated that the silica shell underwent aluminothermic reduction, which led to silicon and an aluminosilicate as intermediate. Thermal analysis indicated that around 90% of Bi element are still in nanoscale.

The average phase-transition temperature is lower than the Bi-Al2O3 nanocomposite prepared by powder metallurgy described in the literature. 2 Based on these results, the next tracks to further explore is the major challenge of reaching phase change nanoparticles suitable for thermal management of Li-ion battery, with an operating temperature at 60-80 °C. This could be achieved by incorporating nanoparticles of low melting point alloys, such as Bi-In alloys.

By this method, we could expect to widen the phase transition temperature range to achieve

. WDXRF analysis of samples with direct incorporation of nanoparticles in molten Al

A master nanocomposite method was proposed by the group of Li via sintering of nanoparticles and Al micropowder (50 wt % loading) together at 900 °C leading to a robust and compact pellet. 5 In our study, this method cannot be directly applied because of the cohesion of the nanoparticles and of the potential reactivity between Al micron scale grains and the particles' boron-rich shell before dispersion. An alternative two-step processing method was then developed. Nanoparticles were dispersed in absolute alcohol in the presence of the aluminum powder. This suspension was dried in an ultrasound bath to ensure its homogeneous deposition over aluminum powder. The mixture was pressed into a pellet, then diluted in an aluminum melt for ultrasound dispersion. This additional step of compaction of the nanopowder with aluminum enhanced the overall incorporation. The macroscale homogeneity of the assynthesized nanocomposites was evaluated by WDXRF. For a HfB2-Al nanocomposite, the Les matériaux à changement de phase ont une application limitée dans le contrôle thermique de la température des batteries des véhicules en raison de leur faible conductivité thermique et de leur faible résistance aux vibrations. 5 L'incorporation de matériaux à changement de phase dans une matrice métallique suscite un grand intérêt dans l'industrie automobile en raison de la haute conductivité thermique, de leurs propriétés mécaniques élevées et de la large gamme de températures de fonctionnement des matrices métalliques. 6 -L'incorporation des particules est réalisée en deux étapes afin d'assurer un transfer quantitatif des particules dans la matrice métallique : 1) mélanger les nanoparticules et l'aluminium à l'état de poudres puis presser ce pré-mélange en pastille, 2) diluer la pastille dans l'Al fondu.

-Un creuset en alumine a été spécialement conçu pour augmenter la hauteur du métal fondu et la résistance au choc thermique.

-Une canne à ultrasons en niobium a été conçue par le fournisseur afin d'éviter la pollution des matériaux finaux par la sonotrode usuelle (alliage de titane).

-Un intervalle on/off du traitement d'ultrason a été appliqué afin d'améliorer la dispersion et de refroidir la sonotrode. L'analyse de la taille moyenne des grains d'aluminium révèle un raffinement de la microstructure, avec des grains de plus petite taille que les grains initiaux avant chauffage. Des échantillons de HfB2/Al à 0,5 vol.% ont été refroidis à différentes vitesses de refroidissement.

La taille moyenne des grains est de 120 µm, 78 µm et 20 µm respectivement pour les taux de refroidissement de 6, 18 et 100°C min -1 . Le refroidissement par azote liquide permet un refroidissement d'environ 100°C min -1 s'avère donc être la meilleure pour affiner la microstructure des MMnCs.

L'effet de renforcement des nanoparticules ont été étudié avec des tests en nanoindentation. La dureté moyenne de l'échantillon avec l'inclusion des nanoparticules HfB2 est de 0,53 GPa, plus élevé que la référence de l'aluminium pur préparée dans les mêmes conditions à 0,42 GPa. Plus de 50% des mesures montrent une dureté plus élevée que l'aluminium pure, correspondant la zone renforcée entre les joints de grains. La moyenne de cette population est de 0,66 GPa, nettement supérieur à l'aluminium pur. Ces mesures confirment l'effet renfort des nanoparticules HfB2.

La cartographie des propriétés mécaniques corrélée avec l'image optique permet de mieux visualiser la différence entre l'intérieur des grains et les joint de grains (figure 2). On La première étape de la préparation du composite consiste en la synthèse de nanoparticules de bismuth avec une taille contrôlée (Figure 3a). Celle-ci a été réalisée selon une procédure de chimie colloïdale précédemment rapportée par Scheele et al. 10 Des nanoparticules de bismuth d'environ 25 nm avec une distribution de taille relativement étroite ont été synthétisées.

Dans un deuxième temps, les nanoparticules de bismuth ont été recouvertes d'une coquille de silice de 7 nm d'épaisseur par la méthode de Stöber. 11 Les nanoparticules de Bi sont encapsulées dans les coquilles et forment des amas de plusieurs particules. Les agrégats ont une taille allant jusqu'à 160 nm (Figure 3b).

Des nanocomposites d'aluminium dans lesquels sont incorporées des nanoparticules de

Bi@SiO2 coeur-coquille (Bi@SiO2-Al) et de Bi (Bi-Al) ont été obtenus par le procédé décrit

Nanoparticules dans des matrices métalliques pour des applications en automobiles Résumé

Ce travail porte sur l'élaboration de nanocomposites à matrices métalliques pour des applications automobiles. Le procédé de fabrication repose sur la dispersion de nanoparticules ayant des capacités de renforcement ou d'absorption de chaleur dans un métal par métallurgie liquide. Après une étude bibliographique pour situer le sujet, le manuscrit porte d'abord sur l'élaboration des nanoparticules comme briques élémentaires, en particulier des particules coeur-coquille Bi@SiO2 à bas point de fusion. Ensuite, plusieurs méthodes sont étudiées pour trouver les meilleures conditions d'incorporation des particules, qui reposent sur le couplage entre le chauffage à induction et le traitement ultrasons dans la matrice fondue. Des nanocomposites combinant des matrices à base d'aluminium et des nanoparticules de SiC ou de HfB2 sont ainsi fabriqués pour renforcer les propriétés mécaniques. La microscopie électronique révèle la ségrégation des particules aux joints de grains. Les mesures de nanoindentation et de microindentation permettent de quantifier la dureté à l'échelle microscopique, révélant ainsi une hétérogénéité des propriétés mécaniques entre les grains et les joints de grains, plus durs en raison de leur forte concentration en particules. Enfin, des nanocomposites métalliques à changement de phase sont élaborés en dispersant dans l'aluminium les particules Bi@SiO2 préalablement synthétisées. L'étude par microscopie électronique de l'interface particules/matrice met en évidence les mécanismes d'incorporation. La coquille de silice protège partiellement les particules de bismuth de la coalescence lors du traitement ultrasons, permettant un abaissement de leur point de fusion. Grâce à cette méthode, nous pourrions élargir la plage de température de transition pour obtenir une absorption progressive de la chaleur avec l'augmentation de la température des batteries. Mots clés : [Nanocomposites, matériaux à changement de phase, propriétés thermiques]

Aluminum matrix nanocomposites for applications in automobile industry Abstract

This PhD work is dedicated to the synthesis of metal matrix nanocomposites for automotive applications, by dispersing functional nanoparticles in an aluminum matrix by liquid metallurgy to improve mechanical and thermal regulation properties. After a bibliographic study to set the context, we describe the preparation of the initial nanoparticles as building blocks, especially core-shell Bi@SiO2 nanoparticles with low melting point. Then, several dispersion methods are explored to find the best conditions for the particle's incorporation, which relies on ultrasonication in the molten matrix during induction heating. We then prepare nanocomposites combining Al-based matrices with SiC or HfB2 nanoparticles, to improve mechanical properties. Electron microscopy study reveals segregation of the nanoparticles at grain boundaries. Nanoindentation and microindentation measurements enable the quantification of the hardness at the microscopic scale, showing a heterogeneity of mechanical properties, with higher hardness at the grain boundaries that are concentrated in nanoparticles. We then explore the fabrication of phase change nanocomposites by incorporating in aluminium the above described Bi@SiO2 particles. An electron microscopy study of the Al/nanoparticle interface sheds light on reaction mechanisms during the incorporation. The silica shell partially protects the nanoparticles against coalescence during the ultrasonic treatment in molten aluminum. We then assessed the role of the inclusion size on the melting point of Bi, evidencing a decrease of the melting point with the particle size. By this method, we could then widen the phase transition temperature range to achieve progressive heat absorption with the increase of the battery temperature. Key words: [Nanocomposites, phase change materials, thermal properties]

  fabricated a bulk reduced graphene oxide-reinforced Al matrix composites with nanolaminated microstructure. Compared with the unreinforced Al matrix, these composites showed significantly improved stiffness and tensile strength, with good yield strength (302 MPa) and elastic modulus (87 GPa) (Figure 1.2). It should be noticed that for all carbon allotropes in Al matrix, a thin layer of Al4C3 could be formed by interfacial reaction, at the expense of the reinforcement objects. It then degrades the reinforcement objects but promotes load transfer. 40 b) Oxides Oxides like Al2O3, ZrO2, TiO2, ThO2 are often added to metallic matrices to provide ductility and fracture toughness from the metallic matrix, creep resistance and high strength from the additives. 11 c) Nitrides Nitrides like nanoparticles of Si3N4 42 and microparticles of Fe3N 43 have been investigated to improve mechanical properties. Recent work focused on layered boron nitride. BN nanotubes embedded in an Al matrix possess excellent elastic modulus (750-1200GPa), strength (61 GPa)

  produced a BN-Al composite, with in situ formation of AlB2 and AlN at the interface. Improved tensile strength of 170 MPa at 500 °C was achieved compared to a value of only 33 MPa for a pure Al sample produced by spark plasma sintering. A nanocomposite reinforced with large-sized cubic BN particles has also been fabricated by cold spray and shows improved abrasive wear resistance. 46 d) Carbides

  firstly prepared well mixed Al powder and 10 vol% SiC nanoparticles via ball milling, then hot pressure sintering and hot extrusion were applied sequentially to get an Al/SiC nanocomposite that showed ultrafine Al grains and intragranular SiC nanoparticles. Compared to pure Al alloy, the nanocomposite exhibited superior mechanical behavior as shown from the stress-strain curve (Figure 1.3). The yield strength and tensile strength are improved by 284% and 259%, respectively, compared to the pure Al alloy.

Figure 1 . 3 .

 13 Figure 1.3. Stress-strain curves of a pure Al alloy and a 10 vol% SiC/Al alloy nanocomposite, (b) Fractograph of Al alloy, (c) Fractograph of the nanocomposite.[START_REF] Li | Enhanced combination of strength and ductility in ultrafine-grained aluminum composites reinforced with high content intragranular nanoparticles[END_REF]

  wettability of reinforcement objects by the matrix. Pourhosseini et al. 61 successfully modified Al2O3 nanoparticles via electroless deposition of Cu, Ni and Co and subsequently applied these particles to reinforce A356 alloy by stir casting. The wettability of the coated nanoparticles was enhanced, resulting in homogeneous distribution in the aluminum matrix. The final mechanical properties (such as microhardness and compressive strength) were then significantly improved, especially for composites encompassing Ni-coated Al2O3. Another method to reach dispersion focuses on the ability of the matrix to carry reinforcement objects. Mousavian et al. 62 managed to reinforce an A356 aluminum matrix with SiC ceramic nanoparticles by combining techniques of stir casting and semi-solid extrusion. The authors found that the addition of nickel acted provided a SiC carrier, leading to a uniform dispersion of particles into the aluminum matrix. In view of the mechanical property, the ultimate tensile strength and elongation of the final Al/SiC composite in presence of nickel could reach around 304 MPa and 5.8%, respectively, while the introduction of titanium decreased the elongation ratio. b) Reinforcement particle size

Figure 1 . 4

 14 Figure 1.4 Energy volumetric and mass densities of different battery systems. 71

Figure 1 . 5

 15 Figure 1.5 Scheme of a traditional Li-ion battery. 77

  , B, I, U and V are the heat generation rate, resistance loss in the cell, reversible entropic heat of the electrochemical reaction, electric current traveling through the cell, open-circuit voltage and the cell voltage of the Li-ion batteries, respectively.

Figure 1 . 7

 17 Figure 1.7 Scheme of a spark plasma sintering device. 102

Figure 1 . 8 .

 18 Figure 1.8. Schematic diagram of ultrasonic assisted casting.103

Figure 1 . 9 . 1 . 4 . 4 techniques 1 . 4 . 4 . 1

 191441441 Figure 1.9. Flow chart of the fabrication of squeeze casted A356/SiC nanocomposites. 60

  Other related works were published recent years, especially from the group of Li 116 who worked on the incorporation of molten-salt synthesized TiB2 and ZrB2 nanoparticles in aluminum. A good dispersion of nanoparticles is essential for the mechanical properties and is a clear bottleneck in the facile fabrication of MMNCs with various reinforcement phases. It requires thorough control of compositional parameters (wettability of the matrix, nature of the compounds) and of processing parameters (heating process, ultrasound treatment power, heating/cooling rate). This part of the work will be explored in chapter 3. Phase change materials have limited applications in vehicle's battery thermal control because of their low thermal conductivity and low resistance against vibration. Composites made by the incorporation of phase change materials in metal matrices are an interesting candidate for the automobile industry due to their high thermal conductivity, good mechanical properties and wide working temperature range. Conventional strategies for MMNCs design include powder metallurgy and bottom-up fabrication, which require high fabrication costs.

  Abbreviations:

Figure 2 . 1 .

 21 Figure 2.1. (a) Powder XRD pattern and (b-d) TEM images of Bi nanoparticles synthesized (b) at 60 °C for 24h from room-temperature oleylamine injection, (c) at 60 °C for 6h from roomtemperature oleylamine injection, the low contrast part around nanoparticles are the organic residue. (d) at 60 °C for 24h from 55 °C oleylamine injection.

Figure 2 . 2 .

 22 Figure 2.2. (a, b) TEM images of Bi@SiO2 nanoparticles with silica shells prepared by the Stöber method. The Bi particles in (a) were washed 2 times before encapsulation. The corresponding size distribution is shown in (c). The Bi particles in (b) were washed 5 times before encapsulation.

Figure 2 .

 2 3 shows the similarity of the dispersibility of samples with or without MUA. Therefore, MUA does not appear as an efficient molecule to trigger ligand exchange and/or dispersion of the Bi particles.

Figure 2 . 3 .

 23 Figure 2.3. TEM images of Bi@SiO2 nanoparticles from Bi nanoparticles washed during 2 cycles before further silica coating: (a) without ligand exchange; (b,c) ligand exchange with 1.5 MUA nm -2 (d) ligand exchange with 3 MUA nm -2 (e) ligand exchange with 5 MUA nm -2 .

Figure 2 . 4

 24 Figure 2.4 Influence of ligand exchange on the dispersibility of bismuth nanoparticles. (a-c) TEM images of Bi nanoparticles after two washing cycles and redispersion in ethanol: (a) without ligand exchange; after ligand exchange with (b) MUA and (c) MHA. (d-f) TEM images of corresponding nanoparticles shown in (a-c), centrifuged at 2000 rpm and redispersed in hexane: (d) without ligand exchange; after ligand exchange with (e) MUA and (f) MHA before redispersion. Low contrast layers around nanoparticles correspond to the organic residuals from the synthesis.

Figure 2 . 5

 25 Figure 2.5 TEM images of Bi@SiO2 samples for initial Bi nanoparticles after two washing cycles and coated with silica (a) without ligand exchange, with ligand exchange by (b) MUA and (c) MHA before silica encapsulation.

Figure 3 . 1

 31 Figure 3.1 TEM images (left) and powder XRD patterns (right) of (a,b) HfB2 nanoparticles coated with an amorphous boron-rich layer, as shown between two arrows; (c,d) of commercial SiC nanoparticles, respectively.

Figure 3 . 2

 32 Figure 3.2 Experimental setup for incorporating nanoparticles into a molten metal by coupling induction heating and high temperature ultrasonication. a) reaction system, b) sonicator with Nb probe, c) alumina crucible.

Figures 3 .

 3 Figures 3.3a,b show SEM images of HfB2-Al nanocomposites prepared in a resistive furnace, with ultrasound treatment through a Nb tip for 5 minutes, followed by cooling down in liquid nitrogen. Aggregation of HfB2 particles was observed along the Al grain boundaries of the final metal matrix rather than a uniform distribution within the matrix. EDS analysis measurements (Figure3.3c and d) yield a heterogeneous composition between the boundaries and the grains, which indicated that HfB2 nanoparticles were segregated towards grain boundaries.

Figure 3 . 3 .

 33 Figure 3.3. Characterization of a 0.8 vol% HfB2-Al nanocomposite prepared with ultrasound treatment by Nb tip for 5 minutes and cooled down in the liquid nitrogen. (a)Optical image and size distribution of Al grains. SEM image (b) and EDS analysis of a grain boundary (c, corresponding to the point of spectrum 1 in b) and inside the Al grain (d, corresponding to the point of spectrum 2 in b). Dark and grey contrasts correspond to aluminum and to aggregates of nanoparticles, respectively.

Figure 3 . 4 .

 34 Figure 3.4. (a,b) TEM images of a 0.8 vol% HfB2-Al nanocomposite prepared by 5 min sonication in a resistive furnace followed by liquid N2 quenching. Dark and light parts correspond to HfB2 nanoparticles and to the aluminum matrix, respectively. (c-e) STEM-EDS mapping showing the overlap of Hf (d) and (e) B elements in the aggregates

Figure 3 . 5 .HfB 2 -

 352 Figure 3.5. Microscopic images of a 0.8 vol% SiC-Al nanocomposite prepared with ultrasound treatment by Nb tip for 5 minutes and cooled down in liquid nitrogen. (a) SEM micrograph, (bd) SEM-EDS mapping, (e) Optical microscopy image of the composite (inset shows the grain size distribution), (f) TEM image of a SiC nanocrystal in the composite.

Figure 3 . 6 . 3 . 3 . 4

 36334 Figure 3.6. Microscopy observations for a 0.8 vol% HfB2-AlSi7Mg0.3Cu0.5 nanocomposite prepared with ultrasound treatment by Nb tip for 5 minutes and cooled down in liquid nitrogen. (a) Optical microscopy image and corresponding size distribution of Al grains, (b) SEM micrograph and (c-g) corresponding SEM-EDS maps

Figure 3 .

 3 Figure 3.7. 0.8 vol% HfB2/Al prepared with a cooling rate of (a) 6, (b)18 and (c) 100 °C min -1

Figure 3 . 8 .

 38 Figure 3.8. Optical microscopy images of 2.4 vol% SiC-Al (a) and 0.8 vol% SiC-Al

Figure 3 .

 3 Figure 3.9 (a) Optical microscopy image, corresponding (b) hardness and (c) elastic modulus of a 0.8% HfB2-Al nanocomposite.

  3 and distribution figure 3.10a). These results were in agreement with microhardness tests, where nanocomposites containing the same volume percentage of SiC nanoparticles were harder than that of HfB2. Measurements at a depth of 1 µm were performed 400 times to reach a statistical assessment of mechanical properties. The cumulative distribution curve (figure 3.10b) for the 0.8 vol% SiC-Al composite is above the curve for the 0.8 vol% HfB2-Al sample, indicating that the reinforcement effect of SiC is more pronounced than HfB2 on the hardness of the final nanocomposite. Nonetheless, HfB2 still shows reinforcement of hardness properties compared to bare aluminum (figure 3.10b).

Figure 3 . 10 .

 310 Figure 3.10. Statistic tests for hardness measured at (a) 3 µm and (b) 1 µm for reference Al, 0.8 vol.% HfB2-Al and 0.8 vol.% SiC-Al

Figure 3 . 11 .

 311 Figure 3.11. Statistical evaluation of hardness measured at 3 µm (a) and 1 µm (b) for 0.8 vol.% SiC-Al, 2.4 vol.% SiC and reference Al

Figure 3 . 13 .Figure 3 . 14 .

 313314 Figure 3.13. Elastic modulus measured at 1 µm

3* 3

 3 *0.5 mm cut in the middle of the sample. Polishing was then performed with a diamond disk to get a 200 μm thick sample. The thickness of the centre of the slab was further reduced to 50 μm by using a precision dimpling instrument (Model 515, South Bay Technology). Ion milling was performed for the final thinning. TEM images were recorded on a Tecnai spirit G2 microscope operating at 120 kV. The STEM images and STEM-EDX mapping were recorded with a JEM 2100Plus UHR microscope and a Tecnai F20 that were both operating at 200 kV.The size distribution of precursor nanoparticles was determined by counting 150 to 200 nanoparticles from TEM images. The size distribution of nanoparticles in the final nanocomposites was determined by counting 70 to 120 particles, from SEM images and TEM images, respectively.

Figure 4 .

 4 2a). The TEM images (Figure 4.2b) showed that the sample consists in bismuth nanoparticles around 25 nm with a relatively narrow size distribution.

Figure 4 . 2 .

 42 Figure 4.2. (a) Powder XRD pattern and (b) TEM images of Bi nanoparticles, (c) Bi@SiO2 nanoparticles and (d) Bi@SiO2 heated at 350 °C for 1 h. Insets in (b) and (d) show the corresponding size distributions of Bi core.

  ) and Bi (Bi-Al) nanoparticles were achieved by ultrasound-assisted casting. To ensure full incorporation of the nanoparticles in molten aluminum, aluminum and nanoparticles powders were firstly dispersed in ethanol by sonication. The mixture was then dried and pressed into a pellet with Bi loading of 0.25 at.%. The particles were then dispersed in aluminium at 740 °C by ultrasound treatment in an induction furnace that ensures efficient heating and melting of the aluminum matrix. The resulting samples showed metallic aluminum appearance (Appendix 4.1). The Bi/Al ratio was evaluated by WDXRF on the top, middle and bottom parts of each sample. It is about 0.3 at. % in all pieces, which is consistent with the initial precursor ratio and indicates that the nanoparticles are homogeneously distributed into the whole volume of the samples.SEM images (Figure 4.3a) highlights white areas corresponding to electron-rich bismuth, with a size up to 120 nm. TEM (Figure 4.3b-d) evidences the existence of 10 nm (Figure 4.3b, c) and 60 nm (Figure 4.3d) nanoparticles of bismuth, as demonstrated by high resolution TEM (HRTEM, Figure 4.3c) and selected area electron diffraction (SAED) (Figure 4.3d). Bigger nanoparticles around 120 nm were also observed in TEM-HAADF images presented in Appendix 4.3. The statistical analysis of SEM and TEM images allows the

Figure 4 . 3

 43 Figure 4.3 SEM and TEM characterization of a Bi@SiO2-Al nanocomposite. (a) backscattered electron imaging-mode SEM image, bright and dark areas correspond to Bi and Al regions, respectively. (b)-(d) TEM images. Dark and light contrast correspond to Bi and Al, respectively. Insets in (a) and (b) show the corresponding size distributions. The inset in (d) shows the selected-area electron diffraction (SAED) pattern corresponding to image (d) and indexed along the structure of bismuth.

Figure 4 . 4 .Figure 4 . 5 .

 4445 Figure 4.4. (a) STEM-HAADF, (b) STEM images and corresponding STEM-EDS maps of Bi@SiO2-Al nanocomposite (c-f),

Figure 4 . 7 .

 47 Figure 4.7. HAADF (a) andTEM image(b) of a Bi-Al nanocomposite and size distribution.

Figure 4 . 9 .

 49 Figure 4.9. DSC analysis of a Bi@SiO2-Al nanocomposite (a) with zoom in the area 250-290 °C during heating (b). for the 1 st and 15 th cycles. The DSC analysis of a reference Bi-Al sample is also shown for the first cycle in (b) (orange).

Figure 1 .

 1 Figure 1. a) Clichés TEM des nanoparticules de HfB 2 avant incorporation . b) Diagramme de DRX de la poudre constituée de nanoparticules HfB 2 . c-d) Clichés TEM du nanocomposite 0,5% HfB 2 -Al. e) Clichés de SEM du nanocomposite 0,5% HfB 2 -Al. e) Clichés SEM du nanocomposite 0,5% SiC-Al.

  observe une dureté hétérogène : une population avec une dureté relativement homogène correspondant aux grains et une deuxième population sensiblement plus élevée pour les zones intergranulaires. Ce phénomène est aussi observé dans d'autres nanocomposites tel que SiC-Al, HfB2-AlSi7Cu0.5Mg0.3.

Figure 2 .Partie 2 .

 22 Figure 2. Cliché de microscopie optique (gauche) correlé à la cartographie des propriétés mécaniques (dureté, module d'élasticité) du nanocomposite 0,5% HfB 2 -Al.

précédemment en partie 1 .

 1 Les images SEM (Figure3c) montrent que les zones riches en bismuth fortement contrastées ont une taille allant jusqu'à environ 100 nm. Les observations TEM (Figure4d) met en évidence l'existence de nanoparticules de bismuth d'environ 10 nm (Figure4.3d) et 60 nm. L'élargissement de la distribution de taille des particules après incorporation, vers les tailles plus petites et plus grosses peut être expliquée par le mécanisme de maturation d'Ostwald au cours du processus d'incorporation, qui vient de la dissolution partielle du bismuth dans l'aluminium fondu, comme attendu d'après le diagramme de phase Bi-Al.12 Nous avons pu mettre en évidence la réduction in situ de la coque de silice par aluminoréduction, qui conduit à du silicium et à un aluminosilicate. Pendant le chauffage assisté par un traitement aux ultrasons, la coquille de silice est séparée du noyau de Bi, cela conduit à la lixiviation des gouttelettes de Bi qui ont subi une maturation d'Ostwald et ont évolué en deux populations de taille.Le nanocomposite Bi-Al contient des particules de Bi de plusieurs micromètres. Ce résultat indique que bien que la coquille de silice ait subi une réaction chimique et une mobilisation locale, elle a agi efficacement contre la coalescence des coeurs par rapport à l'échantillon sans coquille de silice. basée sur la composition chimique (0,47 J g -1 ). les courbes DSC sont restées inchangées après 15 cycles de chauffage-refroidissement entre 20 et 300 °C, ce qui montre la bonne stabilité au cyclage du matériaux.Le nanocomposite Bi-Al présente trois pics endothermiques à 262, 269 et 271 °C au chauffage, et déplacés vers une température plus élevée par rapport au Bi@SiO2-Al, ce qui implique la coalescence progressive des nanoparticules en particules massives pendant le cyclage sans la coquille de SiO2.Par rapport au composite Bi-Al préparé par métallurgie des poudres dans la littérature,6 la distribution des températures de fusion pour Bi@SiO2-Al préparé par notre procédé de métallurgie liquide est plus large et décalée vers des températures plus basses. Cela démontre une diminution de la taille des particules de Bi obtenue par notre protocole. La petite taille initiale des nanoparticules, l'enveloppe protectrice en silice, et le chauffage rapide par le four en induction sont les points clés pour obtenir des petites inclusions avec une large gamme de température de transition, ce qui est un avantage considérable pour l'application finale.

Figure 4 .

 4 Figure 4. Analyse DSC des nanocomposites Bi@SiO 2 -Al (gauche) avec aggrandissement d'une zone au chauffage (droite) pour les 1er et 15e cycles. Anaylse DSC (droite, courbe orange) pour un composite de reference Bi-Al.

Table 1

 1

.1 above listed some of the commercial nanoparticles for MMNCs. They can be categorized into several types: a) Carbon allotropes Figure 1.2. Strengthening and stiffening efficiencies of reduced graphene oxide in a reduced graphene oxide (RGO)/Al nanolaminated composite compared to other reinforcement materials in Al matrix composites

  •°C -1 , high ultimate tensile strength increased by up to 23.68 %, high impact strength and hardness.

	1.2.2.2 Optimized properties of reinforcement
	a) Reinforcement uniformity into matrix phase

  20,[63][64][65][66][67] Reddy et al.67 fabricated aluminum metal matrix nanocomposites reinforced by different contents of SiC nanoparticles of 55 nm diameter (0, 0.3, 0.5, 1.0 and 1.5 vol%) via microwave sintering and hot extrusion techniques. Increasing the SiC content from 0.3 to 1.5 vol% enhanced the final yield strength from 66 to 114 MPa,

Table 1 . 3

 13

: Sound propagation rate in different metals and water

106

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  and hydrothermal methods as well as in microemulsions. Selected examples of syntheses are summarized in Table2.1. under argon under stirring, yielding Bi particles. After washing, 50 mg of nanocrystalline Bi powder could be obtained. The particle size could be adjusted down to ca. 3 nm by changing the [H2O]/[AOT] ratio, but the samples showed high polydispersity.

	were then mixed One of the most prominent colloidal synthesis pathways to monodispersed Bi
	nanoparticle samples has been proposed by the group of Buhro's group 4,5 . Thermal
	decomposition of Bi[N(SiMe3)2]3, prepared by the reaction between Na[N(SiMe3)2]3 and Bi
	salts, was performed in the presence of a polymer surfactant poly(1-hexadecene-co-1-
	vinylpyrrolidinone) (PHD-co-PVP) and NaN(SiMe3)2 at 190 ℃ for 20 h. This synthesis
	resulted in a dark Bi colloidal dispersion. The particle size was restricted to around 21 and
	24 nm.
	Back in 2000, Foos et al. 5 applied the inverse micelle method to produce Bi
	nanoparticles with a high yield and a diameter below 10 nm. The reduction of BiOClO4 occurs
	in an aqueous solution inside of AOT (dioctyl sulfosuccinate, sodium salt) reverse micelles. A
	second micelle solution containing inverse micelles was composed of NaBH4 as reducing agent
	dissolved in aqueous solution with the same [H2O]/[AOT] ratio. Those two micellar systems

Table 2 .

 2 1 Representative examples of synthesis methods for Bi nanoparticles.4

	Method	Reductant	Stabilizing agent	shape	Particle size range	Reference
	Hydrothermal (80 ℃)	N 2 H 4 •H 2 O	Na oleate	nanospheres	25-40 nm	Wang, Y, 2009. 12
	Hydrothermal (160 ℃)	N 2 H 4 •H 2 O	PSS	nanospheres	60-90 nm	Yang, H., 2013. 13
	Hydrothermal (90 ℃)	N 2 H 4 •H 2 O	PEG	nanobelts	10-50 nm	Ma, D., 2015. 14
	Hydrothermal (90 ℃)	NaH 2 PO 2	Tartaric acid	nanobelts	10-100 nm	Ma, D., 2012. 15
	Hydrothermal (RT)	NaBH 4	starch	nanocrystals	10-20 nm	Xia, F., 2014. 2
	Chemical reduction in DMSO	NaBH 4	Sodium citrate	nanocrystals	3.3 nm	Velasco-Arias, D.,
	(RT)					2012. 16
	Solvothermal (185 ℃)	EG	PVP, DMSO	nanospheres	200-400 nm	Li, J., 2009. 17
	Solvothermal (200-300 ℃)	TEG	PVP	nanospheres	60 nm	Kim, H.J., 2015. 18
	Solvothermal (198 ℃)	EG	PVP, urea	Microspheres	400-800 nm	Chen, G., 2009. 19
				nanospheres
	Microwave-assisted solvothermal	DEG; urea	Sodium citrate	nanospheres	100-600 nm	Wu, J., 2010. 20
	Microwave-assisted solvothermal	EG	EG	Microspheres	45 nm-10 mm	Estrada-Flores, M.,
				nanospheres		2016. 21
	Thermal reduction (140-180 ℃)	-	-	nanospheres	170 nm	Carotenuto, G., 2009. 22
	Colloidal synthesis (315 ℃)	OA	OA	nanospheres	10-70 nm	Wang, Z., 2014. 23
	Colloidal synthesis (180 ℃)	NaN(SiMe 3 ) 2	PHD-co-PVP in DIPB	nanodots	3-115 nm	Wang, F., 2008. 6
	Photochemical synthesis in toluene -	TMSC	amorphous nanoparticles 20 nm	Breitwieser, D., 2015. 24
	Photochemical synthesis in	OA	OA, n-heptane	nanoparticles	10 nm	Brütsch, L., 2017. 25
	propylene carbonate			(unspecified shape)
	One-pot cementation method	Fe (powder)	-	nanowires	60 nm width	Wang, Q., 2007. 26
	(150 ℃)
	One-pot cementation method (RT) Zn (powder)	-	nanotube arrays	3-5 nm width	Yang, B., 2003. 27
					>1 mm length
	Two-pot cementation method (RT) Zn, Al (compact	MD, MA, SA	nanoparticles	10-120 nm	Reverberi, A., 2018. 4
		form)		(unspecified shape)

  Bismuth acetate (1 mmol) was mixed with 1-dodecanethiol(11.1 mL) and heated to 45 °C for 45 min under vacuum. The reaction flask was then purged with N2 and heated to 60 °C. Oleylamine (22.2 mL) was quickly added under stirring and the reaction mixture was maintained at 60 °C for 24 h. Finally, Bi nanoparticles were washed twice by an ethanol/chloroform mixture (20/1 vol./vol.) and then re-dispersed in ethanol.

Table 2 .

 2 2 Molecular structure of MUA and MHA

Table 3 .

 3

1 Aluminum alloy AlSi7Mg0.3Cu0.5 composition

14

Table 3 .

 3 2. Microhardness tests for aluminum nanocomposites

	Sample	Ultrasound	Cooling	Average	Average	Average	standard
		treatment	method	hardness	Al grains	grain	deviation
		time		(HV	(HV 0.05)	boundaries
				0.05)		(HV 0.05)	(HV
							0.05)
	Al	5 minutes	Air, in	24.9	/	/	1.3
			crucible
	0.8vol% HfB 2 -Al	5 minutes	Air, in	40.4	25.7	69.3	18.1
			crucible
	2.4vol% SiC-Al	5 minutes	Air, in	45.9	32.9	78.9	22.2
			crucible
	0.8vol% HfB 2 -Al	5 minutes	Liquid	34.9	29.2	39.4	5.0
			N 2
	2.4vol% SiC-Al	5 minutes	Liquid	43.3	36.1	52.0	5.4
			N 2
	0.8vol% HfB 2 -Al	30 minutes	Liquid	33.0	31.0	34.4	1.0
			N 2
	0.8vol% SiC-Al	30 minutes	Liquid	40.9	33.2	52.7	4.9
			N 2
	Al-	30 minutes	Water	67.2	48.1	80.7	9.8
	AlSi 7 Mg 0.5 Cu 0.5
	0.8vol% HfB 2 -Al-	30 minutes	Water	71.7	51.7	91.6	12.8
	AlSi 7 Mg 0.5 Cu 0.5
	AlSi 7 Mg 0.5 Cu 0.5			97.9	80.7	126.8	16.5
	Al	5 minutes	Liquid
			N 2
	Al	5 minutes	Liquid
	AlSi 7 Mg 0.5 Cu 0.5		N 2
	0.8vol% HfB 2 -Al	5 minutes	Liquid
	AlSi 7 Mg 0.5 Cu 0.5		N 2

*HV 0.05: micro vickers hardness measured with a load of 0.05 kgf.

Table 3 .

 3 3 Hardness from nanoindentation (at a depth of 3 µm).

	Matrix	NPs	UT	Cooling	Hardness (GPa)	Hardness of central
			(min)	method				population (GPa)
					Average	SD	Median Average	SD	Median
	Al (Ref)		5	Liquid N 2	0.43	0.03	0.42	0.43	0.02	0.42
	Al	HfB 2 ,	5	Liquid N 2	0.53	0.07	0.54	0.49	0.06	0.49
		0.8 vol%
		(3 wt%)
	Al	SiC,	30	Liquid N 2	0.52	0.07	0.52	0.52	0.03	0.52
		0.8 vol%
		(1 wt%)
	Al	SiC,	5	Liquid N 2	0.80	0.36	0.67	0.68	0.13	0.66
		2.4 vol%
		(3 wt%)
	Al+		5	Liquid N 2	0.77	0.14	0.71	0.74	0.06	0.72
	AlSi 7 Mg 0.3 Cu 0.5
	Al+	HfB 2 ,	30	Liquid N 2	0.80	0.20	0.74	0.75	0.09	0.74
	AlSi 7 Mg 0.3 Cu 0.5	0.8 vol%
		(3 wt %)
	SD stands for standard deviation
	The average of central population eliminates the 25% lowest values and the 25% highest values

Table 3 .

 3 4. Elastic modulus from nanoindentation (at a depth of 3 µm).

	Matrix	NPs	UT	Cooling	Elastic modulus (GPa)	Elastic modulus of
			(min)	method				central population (GPa)
					Average	SD	Median Average SD Median
	Al (Ref)		5	Liquid N 2	79	2	79	79	0.5	79
	Al	HfB 2 ,	5	Liquid N 2	80	4	80	78	3	79
		0.8 vol%
		(3 wt%)
	Al	SiC,	30	Liquid N 2	79	2	78	76	1	78
		0.8 vol%
		(1 wt%)
	Al	SiC,	5	Liquid N 2	95	10	91	91	4	89
		2.4 vol%
		(3 wt%)
	Al+		5	Liquid N 2	88	3	88	88	1	88
	AlSi 7 Mg 0.3 Cu 0.5
	Al+	HfB 2 ,	30	Liquid N 2	86	4	85	86	1	85
	AlSi 7 Mg 0.3 Cu 0.5	0.8 vol%
		(3 wt %)
	SD stands for standard deviation
	The average of central population eliminates the 25% lowest values and the 25% highest values

Nanocomposites à matrice métallique HfB 2 -Al et ses propriétés mécaniques

  Dans ce cas, le matériau à changement de phase est en général un métal à bas point de fusion. Les stratégies conventionnelles, notamment la métallurgie des poudres, nécessitent des coûts de fabrication élevés pour ce type de composites.7 De plus, la température de travail est également limitée : le changement de phase ne se produit qu'à la température de fusion du métal pur. Pour surmonter ces problèmes, une méthode en phase liquide est proposée : les nanoparticules métalliques (Bi) seront encapsulées dans une enveloppe de SiO2 pour empêcher l'immersion dans la matrice. Ensuite, ces nanoparticules seront incorporées dans la matrice métallique fondue par un procédé assisté par ultrasons. Cette partie du travail sera présentée dans la partie 2.

	Partie 1.

Un montage combinant un four à induction et un ultrasonicateur opérant à haute température a été développé pour la préparation des nanocomposites par métallurgie liquide.

Plusieurs paramètres tels que le protocole d'incorporation des nanoparticules, la nature du creuset, la forme du creuset, la puissance des ultrasons, l'intervalle du traitement ultrasons et la rampe du chauffage ont été étudiés afin d'identifier les meilleures conditions pour la synthèse des nanocomposites :

Chapter 2

Synthesis of low melting point

Bi@SiO 2 nanoparticles

We have developed a two-step synthesis combining colloidal synthesis of Bi and Stöber method. The ligand exchange with two well-known ligands for gold system is not efficient in the case of bismuth particles. The dispersibility after the exchange of ligands is a complex phenomenon, which depends greatly on the experimental conditions. Several mechanisms are in competition between surface binding species like thiols, thiolates, Bi-thiolate complexes.

External parameters such as solvent or atmosphere also influence the exchange, as well as the terminal functional group and the chain length. In the future, simulations could be carried out to further understand surface binding and then identify more suitable ligands. The reverse microemulsion method could also be a promising direction. The aqueous droplets in oil are uniform in size and usually smaller than 100 nm. They could act as nanoreactors and produce narrow-size distributions of core-shell nanoparticles. [START_REF] Hull | Introduction to dislocations[END_REF] Chapter 3

Metal matrix nanocomposites with improved mechanical properties

nanoparticles trapped inside the grains, as observed by TEM and discussed above. The average hardness of HfB2-Al inside the grains has no significant improvement compared with that of pure Al, thus confirming full segregation of HfB2 particles at grain boundaries. At fixed concentration, the average hardness for 0.8 vol% SiC nanocomposites was slightly greater than the hardness of HfB2 nanocomposites. This is contrary to the Vickers hardness of bulk SiC and HfB2, of 25.5 and 30.5 GPa, respectively. 11 The difference might be attributed to the following reasons: a) enhanced hardness of the grains due to SiC incorporation in the grains, b) inverse Hall-Petch occurring occurred for HfB2 particles smaller than 10 nm.

When varying the concentration of SiC from 0.8 to 2.4 vol%, the average hardness of SiC-Al increases from 40 HV to 43 HV. The large standard deviations for these values must however be underlined. This may be explained by macro-segregation observed at 2.4 vol%.

Samples processed with different ultrasound treatment period showed no significant change in hardness for 0.8 and 2.4 vol% SiC-Al samples. This result can be confirmed by their similar microstructure according to the microscopic images.

HfB2 nanoparticles were dispersed in the AlSi7Mg0.3Cu0.5 foundry alloy as well. When the two-step incorporation method was used, the foundry alloy and pure Al were mixed with a ratio of 1:1. As expected, the hardness of the reference sample prepared by this method without nanoparticles was lower than the hardness of the initial alloy (67.2 HV for the mixed alloy and 97.9 HV for the foundry alloy). The sample with nanoparticles shows similar hardness compared to the reference sample, with a hardness of 71.7 HV. Considering the high standard deviation around 10 HV, this difference is not significant. This result is in accordance with microscopy observations, where the Si-based solute of the alloy is responsible for the hardness, rather than incorporated nanoparticles.

Nanoindentation tests

We then sought deeper understanding of structure-property relationships in the nanocomposites by using nanoindentation to assess the mechanical properties of the matrix and of reinforcement nanoparticles. Based on the the microindentation results, 4 samples (tables 3.3 and 3.4) were analyzed to study the influence of the heterogeneous composition on the hardness and the elastic nanoparticles enhances the hardness, again at the grain boundaries (Appendix 3.14), consistent with the observations for the Al matrix composites. 3.4). The mapping of the elastic modulus (figure 3.9) follows the same trend as the hardness mapping. Measurements at 1 µm depth (figure 3.13) reveal that the incorporation of 0.8% nanoparticles of HfB2 and SiC could both improve the average elastic modulus by 13% and 16%, respectively. With the increase of SiC nanoparticles to 2.4%, segregated areas at grain boundaries showed high elastic modulus. On contrary, for measurements at 3 µm, samples containing 0.8% of nanoparticles showed similar elastic modulus with Al reference. The less obvious reinforcement effects at 3 µm could be attributed to the increased analysis volume, where the grain boundary effects are averaged out over the matrix grains. As the nanoindentation measurements (Figure 3.13 and 3.14) only focused on the local modulus, these results presented a high standard deviation and could only serve as a rough estimation of the property of final nanocomposite. Mechanical traction measurements are required to further evaluate the macroscopic behavior of the nanocomposites.

Chapter 4

Metal matrix nanocomposites with improved thermal properties ABSTRACT Metal matrix nanocomposites with low melting point metal inclusions exhibit specific mechanical behavior and unique properties for thermal regulation. In this work, we designed aluminum nanocomposites consisting of phase change nanoparticles dispersed in an aluminum matrix via liquid-state processing techniques. To fabricate these composites, we prepared firstly core-shell Bi@SiO2 nanoparticles by colloidal methods, as described in chapter 2. The nanoparticles were subsequently dispersed in liquid aluminum and the final composites obtained by quenching were characterized by XRD, SEM, TEM and XPS. We demonstrate that the silica shell acted as an efficient barrier against extensive coalescence of particles during the casting process, although the size distribution of bismuth nano-objects widened, which demonstrates mass transfer from one particle to the other. According to heat absorption properties measured by DSC, Bi@SiO2-Al nanocomposites are able to store heat by melting through a wider range of melting temperatures centred at lower temperature than for the composite obtained without silica. These properties are stable upon cycling. The casting pathway based on core-shell metal-silica nanoparticles was then efficient to design a heat absorption material with low fabrication cost, thus paving the way to new thermal regulation materials.

Introduction

Phase change materials (PCMs) display a promising range of applications for the thermal control of electronic devices, since they can absorb/release latent heat from the solid-liquid phase transition. 1 The application for battery packs in electric vehicles is especially promising. Indeed, the temperature of Li-ion batteries should be controlled within a certain temperature range to ensure performance, cycling ability and safety. Phase change materials, compared to air/liquid cooling or heat pipe, use less power and they require less complex thermal system design. 2 Jaguemont et al. 3 have reviewed various PCM systems for battery applications and concluded that PCM could be a low-cost, efficient thermal solution if the conductivity of these materials could be improved. Thus, many studies have been performed by mixing phase change materials with metal matrix, possessing better thermal conductivity compared with common polymer matrices. Metallic materials can also provide better flame-retardant properties than organic phase change materials. For instance, metal (Al/Cu) foam coupled with organic PCM as schematically illustrated in fig. 4.1. The results were compared with a reference sample prepared from bare Bi nanoparticles into molten Al, without a protective shell.

Experimental section

Materials

Bismuth acetate (99.99%), Oleylamine (technical grade, 70%), 1-dodenanethiol (98%), tetraethyl orthosilicate (TEOS, 99.99%), ammonia aqueous solution (28-30%) and aluminum powder (99.99%, 320 mesh) were all purchased from Sigma-Aldrich. Aluminum beads (99.9%, around 3mm diameter) were obtained from Alfa Aesar. The reagents were used as received.

Synthesis of bismuth nanoparticles and silica-coated bismuth nanoparticles

Bismuth nanoparticles were prepared by colloidal synthesis in a mixture of organic solvents based on a previous published protocol and already described in chapter 2. 10 Bismuth acetate (1 mmol) was mixed with 1-dodecanethiol (11.1 mL) and heated to 45 °C for 45 min under vacuum. The reaction flask was purged with N2 and heated to 60 °C. Oleylamine (22.2 mL) at room temperature was quickly added under stirring and the reaction mixture was further heated for 24 h. Finally, Bi nanoparticles were washed two times by ethanol/chloroform 100 (20/1 vol./vol.) and re-dispersed in ethanol. The silica-coated bismuth nanoparticles were prepared by the classic Stöber method as follows 11 : tetraethyl orthosilicate (TEOS) was dissolved in absolute ethanol at a concentration of 0.38 M. 10 μL of TEOS solution and 100 μL of ammonia 30% were subsequently added to 1 mL alcoholic bismuth suspension [Bi] = 6.10 -4 mol.L -1 under vigorous agitation. The mixture was heated at 40 °C for 3 hours followed by washing by two cycles of centrifugation-redispersion in ethanol. The final powder was dried in a vacuum oven at 40 °C for 16 hours. In order to evaluate the thermal stability, 20 mg of Bi@SiO2 nanoparticles were annealed in a Naberthem chamber furnace under air condition for 1 h dwell time with a heating rate of 10 °C/min.

Synthesis of Bi-Al nanocomposites

Bi-Al nanocomposites were prepared by ultrasound-assisted casting of a mixture composed of nanoparticles and Al in an induction oven. Briefly, 2 g of aluminum powders and alcoholic bismuth suspension were mixed by sonication at the boiling point for 1 hour to get stoichiometry ratio of 0.3 at. % between Bi and Al. Then the mixture was dried in a vacuum oven at 40 °C for 1 hour, followed by pressing into a 13 mm pellet under 8 tons for 2 minutes.

Subsequently, the pellet and 0.4 g of pure Al beads were heated into an alumina crucible by an induction oven (CEIA, power cube 90 serie 200) at 740 °C under argon atmosphere. After dwelling for 1 minute, an ultrasound power of 80 watts (QSonica, Q500A-220, USA) was applied through a sonication horn (Nb tip with 4 mm diameter) to disperse the nanoparticles for 5 minutes. The final composite was obtained by cooling down the melt in the crucible. It was then cut with a diamond wire saw for further analyses.

X-ray diffraction

X-ray diffraction was performed with a D8 Bruker diffractometer operating at the CuKα radiation in the Bragg Brentano configuration. The ICSD reference 64709 was used for the indexation of bismuth nanoparticles.

Electron microscopy

The cut pieces of the samples were observed with scanning electron microscopy (SEM, Hitachi SU-70 microscope) equipped with an Oxford X-Max 50 mm 2 EDS detector. For preparation, estimation of the size distributions at two different length scales. The TEM investigation of small nanoparticles displayed a distribution with a mean diameter of ca. 8 nm, while SEM investigation showed a wide distribution with an average of 60 nm. Particles below 10 nm showed spherical shape, while the shape of particles above 30 nm was irregular. Therefore, SEM and TEM evidence a widening of the size distribution from initial 25 nm Bi nanoparticles towards both smaller and larger particles. This behaviour is indicative the Ostwald ripening mechanism during the incorporation process, which indicates partial dissolution of bismuth into molten aluminum as expected from the Bi-Al binary phase diagram, which indicate a ca.

1.5 wt.% solubility of Bi in Al at 700 °C. 13 Note that sonication may also exacerbate Ostwald ripening, as already described on polymer composites. 14,15 In addition to the bismuth particles dispersed in the aluminum grains, we have identified thin Bi regions around 1 nm (Appendix 4.4 e,g) between aluminum grains and conformal to grain boundaries. We attribute this feature to bismuth segregation at the Al grains' solidification front during cooling down.

No silicon was detected by using STEM-EDS analysis on a single bismuth nanoparticle of diameter below 10 nm (Appendix 4.4f). Therefore, bismuth and/or silicon could migrate away from each other in molten aluminum. The reduction process of SiO2 (about 30 μm) by molten aluminum has been studied by Wang et al. 18 In the Al-SiO2 system, the reduction of silica by aluminum is thermodynamically favourable (reaction 1 below). Once alumina is formed, it can react with silica and then form mullite or other aluminosilicates (reaction 2), evolving finally towards alumina and silicon (reaction 3).

Wang et al. 18 also demonstrated that after 30 min of reaction at 700 °C after a 5 °C•min - 1 heating rate, a 5 μm-thick shell of alumina was found surrounding unreacted silica particles, while silicon was within the aluminum matrix and formed rod-like Si inclusions. In our study, the silica component consisted in silica shell with a thickness of 7 nm, three orders of magnitude lower that the silica particle size investigated by Wang et al., 18 which supports full reaction of the silica shell at a similar temperature. Nonetheless, aluminothermic reduction in our case is not complete as Si(IV) is not fully reduced in Si(0). On the opposite, the reaction is stopped at an intermediate stage, thus delivering an intermediate aluminosilicate. This behaviour can be attributed to shorter reaction times, as our experiments rely on induction heating at a rate of 350 °C•min -1 and reaction time of 10 min. We achieved sillimanite as intermediate phase.

Generally, sillimanite forms above 776 °C, 19 40 °C over the experimental heating temperature.

DSC analysis for nanocomposites with phase change inclusions

The DSC curves for the composites (Figure 4.9a) also show endothermic peaks during heating and an exothermic peak during cooling indicating solidification of undercooled bismuth. The high temperature regions of DSC curves (Figure 4.9b) show that the melting peaks significant differ from those for the initial particles described above. The Bi@SiO2-Al nanocomposite exhibits two endothermic peaks with an onset temperature at 251°C: a broad melting peak at 263 °C corresponding to nanoparticles' melting and a low intensity sharp peak at 271 °C corresponding to bulk bismuth. The shift of these peaks to higher temperatures compared to unreacted nanoparticles confirms that some of the nanoparticles coalesced during the incorporation of the particles in the melt. The enthalpy ratio for the peaks at 263 °C and at 271 °C is 9:1, indicating that around 90% of Bi element are still nanoscaled, assuming that the melting latent heat of nanoscale Bi is close to the value for bulk Bi. The heat absorption ability evaluated by the enthalpy of the melting process was 0.38 J•g -1 , close to the expected value based on the chemical composition at 0.47 J•g -1 and the melting latent heat fo bulk Bi. This value is 100 times lower compared to Bi@Al2O3-Al prepared by powder metallurgy, 23 but we may improve it by increasing the amount of inclusion particles within the matrix. DSC over 15 consecutive heating-cooling cycles between 20 and 300 °C shows the cycling ability of the Bi core after melting, as the DSC curves remained unchanged during cycling.

General conclusions and perspectives

progressive heat absorption with the increase of the battery's temperature. The next step would then be to select low-cost elements and materials that would render the approach viable for an industrial process. These proofs of concept nonetheless appear to be very encouraging steps towards the development of new metal matrix nanocomposites. Thanks to its adaptability, the described process could be used for boride-based light-weight materials or for thermal management materials for electronic/electrical device through the choice of appropriate inclusions and metal matrix.

Appendix S3.1 Design of the process for the preparation of nanocomposites

S3.1.1 Incorporation methods

The nanoparticles were first incorporated to the molten metal following the same procedure as described in the literature, 1 that is to say direct incorporation of the powders. Attempts have been made by using BN nanosheets, 31 SiC and HfB2 nanoparticles, and Bi2O3 microparticles (several micrometres). Despite the removal of the native aluminum oxide layer at the surface of the metal matrix grains prior to melting, we were not able to incorporate all the nanoparticles inside the matrix by simple stirring methods (Appendix 3.1). For samples dispersed by 5 minutes ultrasound treatment, a loss of more than 90% of initial particles was detected by WDXRF analysis (Appendix 3.2). Such a low inclusion content compared to the initial ratio could be due to the precipitation of nanoparticles since brown powders were observed at the bottom of the crucible. The difficulties to incorporate the nanoparticles into the melt lie in the poor wettability and the high volume fraction between the nanoparticles powder and liquid aluminum. The maximum load of inclusions prepared by this method was evaluated by increasing the nanoparticle ratio in the pellet. The sedimentation of particles at the bottom part of the matrix was observed in the case of HfB2-Al nanocomposites with nanoparticle content over 2 vol.% (Appendix 3.3c). Furthermore, when the loading of inclusion was over 5 vol.%, either for SiC or for HfB2, it was impossible to press nanoparticles/Al powder mixtures into a cohesive pellet under our experimental conditions (8 tons for 2 minutes). A loose and fragile pellet was obtained. Therefore, we chose to study nanocomposites with volume fraction < 2.4 vol.%. It could also be noticed that at this concentration, it is impossible to detect nanoparticles by XRD.

S3.1.2 Improvement of the ultrasonic treatment setup

High intensity ultrasonic waves generate cavitation and acoustic streaming, which is most likely to refine the microstructure, break aggregates and homogenize the distribution of nanoparticles.

However, as mentioned in chapter 1, this method requires a good control of processing condition such as ultrasonic power, ultrasonic treatment time, erosion state of the tip and the

Piece code

Hf/Al at. ratio Expected 0.025 Original 0.02 A1 0.02 B2 0.02 Bottom 0.02 depth of dipping of the tip in liquid metal. Here we designed an experimental set-up adapted to small volume of aluminum melt around 2 mL to optimize these conditions.

Concerning the choice of the tip, most of the commercial microtips are made of titanium alloy.

In our study, titanium tip brought around 1 wt. % of Ti after 10 minutes of ultrasound treatment in pure aluminum. Besides, titanium is a strong boride former that could react with the boron shell of HfB2 nanoparticles, leading to an increased ratio of Ti in WDXRF analysis. Komarov

Appendix 3.4 EDS analysis of HfB2-Al nanocomposites treated by Nb ultrasonic probe for 5 minutes

The form of the crucible was specially designed as well to increase the dipping depth of the ultrasonic tip (3 mm). Compared to the extremely large batches (ca. 1 L) described in previous studies, this set-up is suitable for preliminary studies without needs to scale up the synthesis of nanoparticles.

Previous works have proven that mechanical properties of nanocomposites obtained by discontinuous ultrasonic treatments were better improved than with continuous ultrasonic treatments for 20 min period. 16 An interval of 5s on/ 5s off was then applied for all experiments.

This interval time could also prevent overheating of the ultrasound generator. Furthermore, it was necessary to assess the influence of the ultrasound treatment power. A power varying from 10 to 40% of the nominal power was applied during the preparation of HfB2-Al samples. Strong ultrasonic power over 40% resulted in the splash of liquid on the upper crucible wall leading to 127 a loss of around 42% of the sample. Intermediate ultrasonic power caused less liquid splash, but pores with size varying from 0.05 to 1 mm were formed at the top part of the sample. The formation of these bubbles could be explained by confinement of the cavitation bubbles during solidification process. According to these observations, the best compromise was to use an ultrasound power of 30% for ultrasonic dispersion, followed by 30 s degassing with a power of 10%. These conditions could act as a model to study the interactions between nanoparticles and matrix environment. These experiments proved the feasibility of in situ fabrication of Bi/Al2O3/Al by a casting method. However, the formation of encapsulated inclusions by aluminoreduction is difficult to be applied to other types of inclusions: the reaction condition needs to be adjusted with the redox potential of the metal oxide and it is difficult to achieve other transition temperature.

However, our concept of direct incorporation of inclusion metal@SiO2 nanoparticles in molten Al presents a promising pathway to widen the nature of inclusions. L'examen de la littérature sur les MMnC actuels ouvre de nombreuses possibilités pour le développement de nouveaux matériaux. En ce qui concerne les propriétés mécaniques des MMnCs, les nanoparticules inorganiques commerciales se sont avérées utiles pour améliorer efficacement la résistance et le module d'Young de la matrice métallique. 3 Les nanoparticules de renforcement doivent être conçues pour satisfaire les mécanismes de renforcement et les conditions de la métallurgie à haute température. Les nanoparticules de borure synthétisées par des sels fondus sont de bons candidats car elles sont de petite taille, légères, dures et stables à haute température. 4 À notre connaissance, ce type de nanomatériaux n'a jamais été incoroporé au sein de nanocomposites. Une bonne dispersion des nanoparticules, essentielle pour les propriétés mécaniques finales peut être obtenue par le contrôle des paramètres de composition (mouillabilité de la matrice, nature des composés) ou par la modification des paramètres de mise en oeuvre (procédé de chauffage, puissance du traitement par ultrasons, vitesse de chauffage/refroidissement). Cette partie du travail sera explorée en partie 1.