In the course of this PhD thesis metastable Cu50Zr50-xTix (0≤ x ≤ 10) and (Cu0.5Zr0.5)100-xAlx (5 ≤ x ≤ 8) alloys were prepared and characterised in terms of phase formation, thermal behaviour, crystallisation kinetics and most importantly in terms of mechanical properties. The addition of Al clearly enhances the glass-forming ability although it does not affect the phase formation. This means that the Cu-Zr-Al system follows the characteristics of the binary Cu-Zr phase diagram, at least for Al additions up to 8 at.%. Conversely, the presence of at least 6 at.% Ti changes the crystallisation sequence of Cu50Zr50-xTix metallic glasses and a metastable C15 CuZrTi Laves phase (Fd-3m) precipitates prior to the equilibrium phases, Cu10Zr7 and CuZr2. A structurally related phase, i.e. the “big cube” phase (Cu4(Zr,Ti)2O, Fd-3m), crystallises in a first step when a significant amount of oxygen, on the order of several thousands of mass-ppm (parts per million), is added. Both phases, the C15 Laves as well as the big cube phase, contain pronounced icosahedral coordination and their formation might be related to an icosahedral-like short-range order of the as-cast glass. However, when the metallic glasses obey the phase formation as established in the binary Cu-Zr phase diagram, the short-range order seems to more closely resemble the coordination of the high-temperature equilibrium phase, B2 CuZr. During the tensile deformation of (Cu0.5Zr0.5)100-xAlx bulk metallic glasses where B2 CuZr nanocrystals precipitate polymorphically in the bulk and some of them undergo twinning, which is due to the shape memory effect inherent in B2 CuZr. Qualitatively, this unique deformation process can be understood in the framework of the potential energy landscape (PEL) model. The shear stress, applied by mechanically loading the material, softens the shear modulus, thus biasing structural rearrangements towards the more stable, crystalline state. One major prerequisite in this process is believed to be a B2-like short-range order of the glass in the as-cast state, which could account for the polymorphic precipitation of the B2 nanocrystals at a comparatively small amount of shear. Diffraction experiments using high-energy X-rays suggest that there might be a correlation between the B2 phase and the glass structure on a length-scale less than 4 Å. Additional corroboration for this finding comes from the fact that the interatomic distances of a Cu50Zr47.5Ti2.5 metallic glass are reduced by cold-rolling. Instead of experiencing shear-induced dilation, the atoms become more closely packed, indicating that the metallic glass is driven towards the more densely packed state associated with the more stable, crystalline state. It is noteworthy, that two Cu-Zr intermetallic compounds were identified to be plastically deformable. Cubic B2 CuZr undergoes a deformation-induced martensitic phase transformation to monoclinic B19’and B33 structures, resulting in transformation-induced plasticity (TRIP effect). On the other hand, tetragonal CuZr2 can also be deformed in compression up to a strain of 15%, yet, exhibiting a dislocation-borne deformation mechanism. The shear-induced nanocrystallisation and twinning seem to be competitive phenomena regarding shear band generation and propagation, which is why very few shear offsets, due to shear banding, can be observed at the surface of the bulk metallic glasses tested in quasistatic tension. The average distance between the crystalline precipitates is on the order of the typical shear band thickness (10 - 50 nm) meaning that an efficient interaction between nanocrystals and shear bands becomes feasible. Macroscopically, these microscopic processes reflect as an appreciable plastic strain combined with work hardening. When the same CuZr-based BMGs are tested in tension at room temperature and at high strain rate (10-2 s-1) there seems to be a “strain rate sensitivity”, which could be related to a crossover of the experimental time-scale and the time-scale of the intrinsic deformation processes (nanocrystallisation, twinning, shear band generation and propagation). However, further work is required to investigate the reasons for the varying slope in the elastic regime. As B2 CuZr is the phase, that competes with vitrification, it precipitates in a glassy matrix if the cooling rate is not sufficient to freeze the structure of the liquid completely. The pronounced work hardening and the plasticity of the B2 phase, which are a result of the deformation-induced martensitic transformation, leave their footprints in the stress-strain curves of these bulk metallic glass matrix composites. The behaviour of the yield strength as a function of the crystalline volume fraction can be captured by the rule of mixtures at low crystalline volume fractions and by the load bearing model at high crystalline volume fractions. In between both of these regions there is a transition caused by percolation (impingement) of the B2 crystals. Furthermore, the fracture strain can be modelled as a function of the crystalline volume fraction by a three-microstructural-element body and the results imply that the interface between B2 crystals and glassy matrix determines the plastic strain of the composites. The combination of shape memory crystals and a glassy matrix leads to a material with a markedly high yield strength and an enhanced plastic strain. In the CuZr-based metastable alloys investigated, there is an intimate relationship between the microstructure and the mechanical properties. The insights gained here should prove useful regarding the optimisation of the mechanical properties of bulk metallic glasses and bulk metallic glass composites.:Abstract/Kurzfassung . . . . . . . . . . . . . . . . . . . . . . . . vii Aims and objectives . . . . . . . . . . . . . . . . . . . . . . . . xiii 1 Metallic glasses and bulk metallic glasses . . . . . . . . . . . . . . . . . . . . . . 1 1.1 Structure of metallic glasses . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 1.2 Glass formation and transformation kinetics . . . . . . . . . . . . . . . . . . 4 1.2.1 Crystallisation kinetics . . . . . . . . . . . . . . . . . . . . . . . . . . . 6 1.2.2 Glass-forming ability . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 1.2.3 Fragility concept of metallic glasses . . . . . . . . . . . . . . . . . . . 10 1.3 Mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13 1.3.1 The potential energy landscape concept . . . . . . . . . . . . . . . . . 16 1.3.2 Role of the shear modulus upon flow of a glass . . . . . . . . . . . . . 20 1.3.3 Factors affecting plastic deformation of BMGs . . . . . . . . . . . . . 25 1.4 Metastable Cu-Zr-based alloys . . . . . . . . . . . . . . . . . . . . . . . . . . 30 1.4.1 Binary Cu-Zr glasses . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32 1.4.2 Minor additions of Al and Ti to glassy Cu-Zr . . . . . . . . . . . . . . 33 2 Synthesis and characterisation methods . . . . . . . . . . 35 2.1 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35 2.1.1 Melt spinning . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36 2.1.2 Cu-mould suction casting . . . . . . . . . . . . . . . . . . . . . . . . . 37 2.2 X-ray diffraction/in-situ experiments . . . . . . . . . . . . . . . . . . . . . . . 38 2.3 Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38 2.3.1 Optical microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38 2.3.2 Scanning electron microscopy . . . . . . . . . . . . . . . . . . . . . . . 39 2.3.3 Transmission electron microscopy . . . . . . . . . . . . . . . . . . . . 39 2.4 Calorimetry/ Dilatometry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39 2.5 Ultrasound velocity measurements . . . . . . . . . . . . . . . . . . . . . . . . 40 2.6 Mechanical testing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 41 3 Effect of oxygen on Cu-Zr-(Ti) alloys . . . . . . . . . . . . . . . . . . . . . . . . 43 3.1 Influence of casting parameters . . . . . . . . . . . . . . . . . . . . . . . . . . 43 3.2 Phase formation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45 4 Effect of Ti and Al on Cu-Zr glasses . . . . . . . . . . . . . . . . . . . . . . . . 53 4.1 Phase formation and thermal stability . . . . . . . . . . . . . . . . . . . . . . 53 4.2 Crystallisation kinetics and fragility . . . . . . . . . . . . . . . . . . . . . . . 64 4.2.1 Isothermal calorimetric measurements . . . . . . . . . . . . . . . . . . 64 4.2.2 Isochronal calorimetric measurements . . . . . . . . . . . . . . . . . . 67 4.3 Structure of Cu-Zr-(Al/Ti) glasses . . . . . . . . . . . . . . . . . . . . . . . . 71 5 Glassy Cu-Zr-(Al/Ti) alloys . . . . . . . . . . . . . . . . . . . . . . . . 79 5.1 Deformation behaviour of glassy ribbons . . . . . . . . . . . . . . . . . . . . 79 5.2 Deformation behaviour of bulk metallic glasses . . . . . . . . . . . . . . . . . 83 5.2.1 Compression tests of Cu50Zr50 . . . . . . . . . . . . . . . . . . . . . . 83 5.2.2 Tensile tests of (Cu0.5Zr0.5)100-xAlx . . . . . . . . . . . . . . . . . . . . 85 5.2.3 Fractography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102 5.2.4 High-strain rate tensile tests . . . . . . . . . . . . . . . . . . . . . . . . 104 6 Cu-Zr intermetallic compounds . . . . . . . . . . . . . . . . . . . . . . . . 111 6.1 Deformation behaviour of Cu10Zr7 and CuZr2 . . . . . . . .. . . . . . . . 111 6.2 Deformation behaviour of B2 CuZr . . . . . . . . . . . . . . . . . . . . . . . . 113 6.3 Relation between intermetallics and BMGs . . . . . . . . . . . . . . . . . . . 119 7 Cu-Zr-(Al/Ti) BMG matrix composites . . . . . . . . . . . . . . . . . . . . . . . . 123 7.1 Microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 123 7.2 Deformation behaviour . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 126 8 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . 137 9 Outlook . . . . . . . . . . . . . . . . . . . . . . . . 139 10 Appendix . . . . . . . . . . . . . . . . . . . . . . . . 143 10.1 Isochronal transformation kinetics (Kissinger) . . . . . . . . . . . . . . . . 143 10.2 Isothermal crystallisation kinetics (Johnson-Mehl-Avrami) . . . . . . . 144 10.3 The fragility concept of metallic glasses . . . . . . . . . . . . . . . . . . . . . 144 10.4 Flow of liquids in the PEL picture . . . . . . . . . . . . . . . . . . . . . . . . . 146 10.5 The interstitialcy theory . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147 Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . 149 Bibliography . . . . . . . . . . . . . . . . . . . . . . . . 151

<p>This dissertation addresses a key 21st Century Grand Challenge – "Combatting Nuclear Terrorism”. A principal component associated with addressing this challenge pertains to timely and near real-time detection and tracking of small quantities of special nuclear materials (SNMs); the isotopes of uranium (U-235), and Plutonium (Pu) which constitute the key components of nuclear warheads. Such detection and tracking, especially for shielded U-235 using passive means is virtually impossible due to the extremely faint neutron-photon emission signals from radioactive decay which can be readily masked. Active photon and/or neutron interrogation methods are the only viable means for HEU detection but the field suffers from detector saturation in extreme 10⁴ R h^-1 radiation fields. Pu isotopes in multi-kg levels emanate neutrons from spontaneous fission that offer a means for passive interrogation with directionality, even at low levels assuming novel, high-efficiency detectors are available. Both U-235 and Pu isotopes also emit Rn gas (an alpha radiation emitter) at trace levels, during decay - which offers a possible novel means for identifying the presence of SNMs – from the faint multi Bq m^-3(pCi L^-1) alpha emitting gas and progeny in air - if only a real time sensitive enough detector were available. </p> <p> </p> <p>This thesis work was aimed at filling critical technology gaps, via researching and advancing the field of metastable fluid detector (TMFD) technology pertaining to novel/transformational passive and active (photoneutron) interrogation of SNMs. The results of R&D from this dissertation provide evidence for rapidly and conclusively monitoring for the presence of Rn-222 and progeny in air at ultra-trace (pCi L^-1) levels – even below the action levels mandated by the U.S. EPA by the development of protocols for sampling and detection using centrifugally tensioned metastable fluid detectors (CTMFD). </p> <p> </p> <p>For SNM neutron emission (either spontaneous or induced) based active and passive interrogation this dissertation presents evidence for advancing into novel designs, and schemes resulting in 100-1000x enhancements in detection efficiency for the acoustically driven ATMFD architecture in single and array forms. Novel drive modes: a direct (fixed and sweep) resonance mode, and radically novel indirect traveling wave mode were used to expand ATMFD capabilities and efficiencies beyond previous iterations of ATMFD technology. The experimentation work has been coupled with multi-physics theoretical modeling and simulations benchmarked against experimental data. ATMFDs in single and array-based architectures are being investigated for offering a novel, high-efficiency means for passive interrogation of SNMs. Coupled together with the Rn-alpha sensing approach, the ATMFD sensors for neutron monitoring enable a first-of-a-kind transformational dual mode architecture for monitoring both HEU (U-235) and Pu based SNMs.</p> <p> </p> <p>Successful results were also demonstrated for rapid and convincing 9 MeV (end point x-ray) photoneutron based active interrogation of 4.5 kg of depleted uranium in ultra-high gamma background of ~10⁴ R h^-1 using a single CTMFD or ATMFD sensor. Under such intense gamma backgrounds, conventional detectors are known to get saturated and have presented a major challenge. The research from this thesis offers a novel solution for both passive and active SNM interrogation. </p>

Ce travail de thèse démontre une technique de détection capable de mesurer, avec une sensibilité à l'atome unique, l'espace des impulsions d'un gaz ultrafroid chargé dans un réseau optique 3D. Nous avons développé un détecteur basé sur des galettes de micro-canaux, capable de sonder électroniquement des nuages d'Hélium-4 métastable. Le gaz est détecté après un temps de vol de 325ms, suffisamment long pour atteindre l'expansion de champ lointain, où la distribution spatiale du gaz coïncide avec la distribution d'impulsion asymptotique. En se plac{c}ant dans un régime proche du remplissage unitaire du réseau, les effets de collisions entre atomes aux premiers instants de l'expansion deviennent négligeables, et donc la distribution d'impulsion asymptotique est égale à la distribution d'impulsion in situ. Nous démontrons expérimentalement cette égalité en comparant nos mesures en champ lointain avec la distribution d'impulsion calculée à partir de l'Hamiltonien de Bose-Hubbard, gr^ace à des simulations Monte Carlo Quantique. Nous observons un bon accord avec la théorie sur plus de 3 ordres de grandeur en densité. Ces simulations sont calculées à partir de nos paramètres expérimentaux, la température étant la seule variable ajustable. Nous utilisons ensuite cette comparaison pour réaliser une thermométrie précise du gaz sur réseau, permettant une exploration de la transition superfluide-gaz normal à travers la mesure directe de différentes quantités, comme la fraction condensée ou la fonction de corrélation à deux particules. / In this thesis, we report the demonstration of a detection technique able to probe, with a single-atom sensitivity, the momentum distribution of an ultracold gas loaded inside a 3D optical lattice. We have developed a micro-channel plate detector, able to electronically probe clouds of metastable Helium-4. The gas is detected after a time-of-flight of 325ms, long enough to reach the far-field expansion, where the spatial distribution of the cloud can be mapped to the asymptotic momentum distribution. By putting ourselves in a regime where the lattice filling is close to unity, the atomic collisions in the first instant of the expansion become negligible, so that the asymptotic momentum distribution is equal to the in situ momentum distribution. We experimentally demonstrate this equality, by comparing our far-field measurements with the momentum distribution calculated from the Bose-Hubbard Hamiltonian, thanks to ab initio quantum Monte Carlo simulations. We show a good agreement with the theory over more than 3 orders of magnitude in density. Those simulations are calculated with our experimental parameters, the temperature being the only adjustable variable. We then use this comparison to perform a precise thermometry of the lattice gas, allowing us to explore the superfluid-normal gas transition through a direct measurement of different quantities, like the condensed fraction or the two-particles correlation function.

Deux expériences d'Optique Atomique Quantique sont étudiées dans ce manuscrit : l'expérience Hong-Ou-Mandel atomique ainsi qu'une expérience préliminaire au test des inégalités de Bell sur la variable d'impulsion de deux atomes. Ces deux expériences permettent de révéler des comportements très particuliers de la matière. La première produit une interférence à deux particules et la seconde suggère l'intrication de deux atomes en impulsion. Ces deux aspects de la mécanique quantique sont non-intuitifs et surprenants. Pouvoir les réaliser avec quelques atomes dans des environnements contrôlés est nécessaire à la consolidation de nos connaissances des lois de la nature et tout particulièrement quand mécanique quantique et gravitation sont à l'oeuvre simultanément. / We report two Quantum Atom Optics experiments : the atomic Hong-Ou-Mandel experiment and preliminary results toward Bell's inequality test on momentum of two atoms. Both experiments reveal interesting behaviors of matter. The first one shows a 2-particle interference and the second one suggests entanglement momenta between two atoms. These aspects of quantum mechanics are really surprising and unintuitive. Realizing these experiments with a few atoms in a controlled environment are necessary to strengthen our understanding of the world, particularly when quantum mechanics and gravitation act simultaneously.

Das hier vorgestellte Vorhaben hat die Entwicklung eines neuartigen Sensor- und Messtechnikkonzepts zur Erfassung der Maximalwerte mechanischer Belastungen von Bauteilen bestehend aus Kohlenstofffaserverbundwerkstoffen mit polymerer Matrix (CFK) zum Ziel. Dazu wurden Untersuchungen zur Anwendbarkeit der spontanen Magnetisierung durch die martensitische Gefügeumwandlung metastabiler Stähle als möglicher Sensoreffekt durchgeführt. Als ein einfaches passives Sensorverfahren hat die Ausbildung einer spontanen Magnetisierung durch die Gefügeumwandlung ein denkbares Anwendungspotenzial. Es ist keine permanente Energieversorgung des Messsystems erforderlich, der Sensor ist vollständig und ohne Kontakt zur Umgebung im Bauteil einbettbar und eine äquivalente Information der maximalen Last ist unauslöschbar im Sensor eingeprägt. Durch kontakt- und zerstörungsfreie Verfahren kann das Auslesen erfolgen. Um die Eignung des Effekts zu überprüfen, erfolgten verschiedene Untersuchungen. Die Ergebnisse von Zugversuchen, durchgeführt an eingebetteten und nicht eingebetteten Sensorproben, und die Messwerte magnetischer Messungen der Sensormagnetisierung wurden zur Ermittlung einer Abhängigkeit zwischen dem Grad der Belastung und der Änderung der Magnetisierung infolge der martensitischen Transformation korreliert. Mittels metallografischer Untersuchungen sowie durch Röntgenbeugung erfolgte zudem der Vergleich zwischen dem Martensitgehalt und den Messergebnissen zur spontanen Magnetisierung. Der vorliegende Artikel beschreibt die Grundlagen des ausgewählten Sensorprinzips, die durchgeführten Arbeiten und erläutert die Ergebnisse der Untersuchungen.

High-entropy alloys (HEAs) have attracted significant interest in the materials science community over the last 15 years. At the first moment, what caught the attention was the fact that these alloys tend to form solid solutions at room temperature, despite being composed of multiple elements in equiatomic or near-equiatomic concentrations. It was initially concluded that the configurational entropy plays a key role in the stabilization of the solid solutions. Later studies revealed the importance of lattice strain enthalpies, enthalpies of mixing, structural mismatch of constituents, and kinetics in phase formation/stability. The study presented in this thesis was branched into three major parts, all related to understanding phase formation, stability, or metastability in this class of alloys. The first part deals with developing an empirical method to predict single-phase solid solution formation in multi-principal element alloys. The second, which makes the core of this thesis, are non-equilibrium solidification studies of CrFeNi and CoCrNi medium-entropy alloys, and CoCrFeNi, Al0.3CoCrFeNi, and NbTiVZr high-entropy alloys. The last part is devoted to understanding the thermophysical properties of CrFeNi, CoCrNi, and CoCrFeNi medium- and high-entropy alloys. An empirical approach, based on the theoretical elastic-strain energy, has been developed to predict the phase formation and its stability for complex concentrated alloys. The conclusiveness of this approach is compared with the traditional empirical rules based on the atomic-size mismatch, enthalpy of mixing, and valence-electron concentration for a database of 235 alloys. The proposed “elastic-strain energy vs. valence-electron concentration” criterion shows an improved ability to distinguish between single-phase solid solutions, mixtures of solid solutions, and intermetallic phases when compared to the available empirical rules used to date. The criterion is especially strong for alloys that precipitate the μ phase. The elastic-strain-energy parameter can be combined with other known parameters, such as those noted above, to establish new criteria which can help in designing novel complex concentrated alloys with the on-demand combination of mechanical properties. The solidification behavior of the CoCrFeNi high-entropy alloy and the ternary CrFeNi and CoCrNi medium-entropy suballoys has been studied in situ using high-speed video-camera and synchrotron X-ray diffraction (XRD) on electromagnetically levitated samples at Leibniz Institute for Solid State and Materials Research Dresden (IFW Dresden) and German Synchrotron DESY, Hamburg. In all alloys, the formation of a primary metastable body-centered cubic bcc phase was observed if the melt was sufficiently undercooled. The delay time for the onset of the nucleation of the stable face-centered cubic fcc phase, occurring within bcc crystals, is inversely proportional to the melt undercooling. The experimental findings agree with the stable and metastable phase equilibria for the (CoCrNi)-Fe section. Crystal-growth velocities for the CrFeNi, CoCrNi, and CoCrFeNi medium- and high-entropy alloys, extracted from the high-speed video sequences in the present study, are comparable to the literature data for Fe-rich Fe-Ni and Fe-Cr-Ni alloys, evidencing the same crystallization kinetics. The effect of melt undercooling on the microstructure of solidified samples is analyzed and discussed in the thesis. To understand the effect of Al addition on the non-equilibrium solidification behavior of the equiatomic CoCrFeNi alloy, the Al0.3CoCrFeNi HEA has been studied. While the quaternary alloy melt could be significantly undercooled, this was not possible in the five-component alloy. Therefore, the investigations on phase formation, crystal growth, and microstructural evolution were confined to the low undercooling regime. In situ XRD measurements revealed that the liquid crystallized into a fcc single-phase solid solution at this undercooling level. However, ex situ XRD revealed the precipitation of the ordered L12 phase for a sample solidified with ΔT = 30 K. Crystal growth velocities are shown to be smaller than in the CoCrFeNi, CrFeNi, and CoCrNi alloys; nonetheless, they are in the same order of magnitude. Spontaneous grain refinement, without the formation of crystal twins, is observed at low undercooling of ΔT = 70 K, which could be explained by the dendrite tip radius dependence on melt undercooling. In situ studies of the equiatomic NbTiVZr refractory high-entropy alloys revealed the effect of processing conditions on the high-temperature phase formation. When the melt was undercooled over 80 K, it crystallized as a bcc single-phase solid solution despite solute partitioning between the dendritic and interdendritic regions. When the sample was solidified from the semisolid state, it resulted in the formation of two additional bcc phases at the interdendritic regions. The crystal growth velocity, as estimated from the high-speed videos, showed pronounced sluggish kinetics: it is 1 to 2 orders of magnitude smaller compared to literature data of other medium and high-entropy alloys. The study of the linear expansion coefficient α and heat capacity at constant pressure 𝐶𝑝 of the equiatomic CoCrFeNi and the medium-entropy CrFeNi and CoCrNi alloys revealed an anomalous behavior with S-shaped curves in the temperature range of 700 – 950 K. The anomalous behavior is shown to be reversible as it occurred during the first and second heating. However, a minimum is only observed on the first heating, while in the second heating a sudden increase of both the α and 𝐶𝑝 occurs at the temperature of the onset of the minima in the first heating. Magnetic moment measurements as a function of temperature showed that the observed anomaly is not associated with the Curie temperature. Consideration of the structural and microstructural evaluation discards a first-order phase transformation or recrystallization as probable causes, at least for the CoCrFeNi and CoCrNi alloys. Based on literature evidence, the anomalies in the temperature dependences of the linear expansion coefficient and heat capacity are believed to be caused by a chemical short-range order transition known as the K-state effect. However, to reveal the exact nature of this phenomenon, further experimental and theoretical studies are required, which is outside the frame of the present work.:Abstract ....................................................................................................................... I Kurzfassung .............................................................................................................. IV Chapter 1: Motivation and Fundamentals .................................................................. 1 1.1 Introduction .......................................................................................................... 1 1.2 The high-entropy alloy (HEA) design concept ...................................................... 4 1.3 Empirical rules of phase formation for HEAs ....................................................... 6 1.4 Calculation of phase diagrams of HEAs ............................................................. 18 1.5 The core effects of HEAs ................................................................................... 20 1.5.1 Lattice distortion .............................................................................................. 20 1.5.2 Sluggish diffusion ............................................................................................ 22 1.5.3 Cocktail effect................................................................................................... 23 1.6 Mechanical properties ........................................................................................ 24 1.6.1 Lightweight high-entropy alloys ....................................................................... 24 1.6.2 Overcoming the strength-ductility tradeoff ...................................................... 26 1.6.3 Cryogenic high-entropy alloys ......................................................................... 28 1.6.4 Refractory high-entropy alloys ........................................................................ 30 1.7 Functional properties .......................................................................................... 33 1.7.1 Soft magnetic properties ................................................................................. 33 1.7.2 Magnetocaloric properties ............................................................................... 35 1.7.3 Hydrogen storage ............................................................................................ 36 Chapter 2: Experimental .......................................................................................... 38 2.1 Sample preparation ............................................................................................ 38 2.2 Electromagnetic levitation .................................................................................. 40 2.3 In situ X-ray diffraction ........................................................................................ 43 2.4 Microstructural and structural analysis ............................................................... 44 2.5 Thermal analysis ................................................................................................ 45 2.6 Dilatometry ......................................................................................................... 45 2.7 Magnetic moment ............................................................................................... 46 2.8 Heat treatment ................................................................................................... 46 Chapter 3: In situ study of non-equilibrium solidification of CoCrFeNi high-entropy alloy and CrFeNi and CoCrNi ternary suballoys ...................................................... 47 3.1 Introduction ........................................................................................................ 47 3.2 Results ............................................................................................................... 48 3.2.1 In situ synchrotron X-ray diffraction ................................................................. 48 3.2.2 High-speed video imaging ............................................................................... 52 3.2.3 Microstructure of the solidified samples .......................................................... 62 3.3 Discussion .......................................................................................................... 64 3.3.1 bcc-fcc nucleation and growth competition ..................................................... 64 3.3.2. Crystal growth kinetics ................................................................................... 68 3.3.3. Microstructural evolution ................................................................................ 70 Chapter 4: The effect of Al addition to the CoCrFeNi alloy on the non-equilibrium solidification behaviour.............................................................................................. 72 4.1 Introduction ........................................................................................................ 72 4.2 Results and Discussion ...................................................................................... 73 Chapter 5: Non-equilibrium solidification of the NbTiVZr refractory high-entropy alloy ................................................................................................................................. 84 5.1 Introduction ........................................................................................................ 84 5.2 Results ............................................................................................................... 85 5.2.1 In situ synchrotron X-ray diffraction ................................................................. 85 5.2.2 Room temperature synchrotron X-ray diffraction ............................................ 88 5.2.3 High-speed video imaging ............................................................................... 89 5.2.4 Microstructure and structure analysis ............................................................. 91 5.3 Discussion .......................................................................................................... 94 5.3.1 Phase formation upon solidification ................................................................ 94 5.3.2 Crystal growth kinetics .................................................................................... 98 5.3.3 Structural and microstructural features............................................................ 99 Chapter 6: Solid-state thermophysical properties of CrFeNi, CoCrNi, and CoCrFeNi medium- and high-entropy alloys ........................................................................... 101 6.1 Introduction ...................................................................................................... 101 6.2 Results ............................................................................................................. 102 6.3 Discussion ........................................................................................................ 106 6.3.1 Thermophysical properties ............................................................................ 106 6.3.2 Short-range order in medium- and high-entropy alloys ................................. 109 Chapter 7: Summary ............................................................................................... 111 7.1 Empirical rule of phase formation of complex concentrated alloys ................... 111 7.2 Non-equilibrium solidification of medium- and high-entropy alloys ................... 111 7.3 Thermophysical properties of the medium- and high-entropy alloys ................ 113 Chapter 8: Outlook ................................................................................................. 115 Appendix 1 .............................................................................................................. 117 Appendix 2 ............................................................................................................. 123 Appendix 3 ............................................................................................................. 133 Appendix 4 ............................................................................................................. 134 References.............................................................................................................. 140 Acknowledgments .................................................................................................. 164 List of publications .................................................................................................. 166 Erklärung ......................................................................................................................... 167

No description available.

Study of the metastable phases obtained by non-equilibrium processing techniques has come a long way during the past five decades. New metastable phases have often given new perspectives to the research on synthesis of novel materials systems. Metastable materials produced by two non-equilibrium processing methods were studied for this dissertation- 304-type austenitic stainless steel (SS304 or Fe-18Cr-8Ni)+aluminum coatings produced by plasma enhanced magnetron sputter-deposition (PEMS) and nanocrystalline Ti, Zr and Hf powders processed by mechanical milling (MM). The objective of the study was to understand the crystallographic and microstructural aspects of these materials. Four SS304+Al coatings with a nominal Al percentages of 0, 4, 7 and 10 wt.% in the coatings were deposited on an SS304 substrate by PEMS using SS304 and Al targets. The as-deposited coatings were characterized by x-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM) and three-dimensional atom probe microscopy (3DAP). Surface morphology and chemical analysis were studied by SEM. Phase identification was carried out by XRD and TEM. The microstructural features of all the coatings, as observed in the TEM, consisted of columnar grains with the columnar grain width (a measure of grain size) increasing with an increase in the Al content. The coatings had grains with average grain sizes of about 100, 290, 320 and 980 nm, respectively for 0, 4, 7 and 10 wt.% Al. The observed grain structures and increase in grain size were related to substrate temperature during deposition. XRD results indicated that the Al-free coating consisted of the non-equilibrium ferrite and sigma phases. In the 4Al, 7Al and 10Al coatings, equilibrium ferrite and B2 phases were observed but no sigma phase was found. In 10Al coating, we were able to demonstrate experimentally using 3DAP studies that NiAl phase formation is preferred over the FeAl phase at nano scale. During mechanical milling of the hexagonal close packed (HCP) metals Hf, Ti and Zr powders, unknown nanocrystalline phases with face centered cubic (FCC) structure were found. The FCC phases could be either allotropes of the respective metals or impurity stabilized phases. However, upon MM under high purity conditions, it was revealed that the FCC phases were impurity stabilized. The decrease in crystallite size down to nanometer levels, an increase in atomic volume, lattice strain, and possible contamination were the factors responsible for the transformation.

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Ce travail porte sur l’étude et la réalisation d’une nouvelle technique de décélération d’un jet supersonique de particules paramagnétiques en utilisant une onde de champ magnétique progressive co-mobile. Cette technique repose sur une méthode de ralentissement basée sur les forces de type Stern Gerlach agissant sur un système paramagnétique en mouvement en présence d’un champ magnétique co-propageant. Cette méthode très innovatrice a l’avantage de pouvoir s’appliquer à une grande palette d’espèces ouvrant ainsi de nouvelles possibilités d’applications. On décrit une approche théorique adaptée qui permet de faire un lien direct entre la théorie, la programmation des paramètres expérimentaux, les résultats obtenus et ce d’une manière systématique, rationnelle et prédictive.Ce mémoire est composé de trois parties. La première porte sur les forces décélératrices et le calcul des différentes forces, de type Stern Gerlach, utilisées dans nos expériences. Les formules établies dans cette partie sont essentielles pour l’interprétation des résultats expérimentaux. La deuxième partie porte sur le dispositif expérimental : le jet supersonique pré-refroidi, la zone d’interaction et la détection. On donne le détail de la réalisation des circuits créant les champs magnétiques nécessaires au guidage et à la décélération du jet. La troisième partie porte sur les résultats des expériences réalisées et leur interprétation directement à partir des équations du mouvement de l’effet Stern Gerlach. Des simulations sont présentées pour étayer les interprétations. On présente les résultats de décélération obtenus récemment sur l’argon et le néon métastables. Ces résultats valident clairement l’importance de l’ajout d’un champ magnétique uniforme qui définit un axe de quantification adiabatique global pour toutes les particules du jet et permet le découplage entre la précession des moments magnétiques et l’action des forces de gradient. Ces résultats mettent en évidence, aussi, l’effet de polarisation du jet qui dépend du sens relatif du champ magnétique uniforme ajouté par rapport à l’onde de champ magnétique progressive.Enfin, la compréhension et le contrôle de la dynamique du piégeage à une vitesse donnée, de l’accélération et de la décélération nécessitent le découplage entre les effets transverses et les effets longitudinaux de l’onde. Ces derniers sont clairement visibles quand le champ magnétique uniforme ajouté vient limiter les effets transverses de l’onde de champ magnétiques progressive. Les perspectives pour ce nouveau décélérateur Zeeman Stern Gerlach sont grandes. Un premier résultat de piégeage du di-azote métastable à 560m/s est présenté et ceci ouvre la voie pour décélérer les molécules paramagnétiques en jet supersonique pulsé. La décélération des radicaux libres et des neutrons est aussi envisageable. / This work focuses on the study and implementation of a new technique of deceleration of a supersonic beam of paramagnetic particles using a co-moving progressive wave of magnetic field. This technique relies on a method of slowing based on Stern-Gerlach forces acting on a paramagnetic system in motion in the presence of a co-propagating magnetic field. This highly innovative approach has the advantage of being applicable to a wide range of species and opens up new opportunities. A suitable theoretical approach is followed, that allows for a direct link between theory, programming of experimental parameters, and experimental results in a systematic, rational and predictive manner.This thesis is composed of three parts. The first concerns the calculation of the various Stern Gerlach forces used in our experiments to decelerate the paramagnetic particles. Formulas established in this section are essential for the interpretation of experimental results. The second part is devoted to the experimental device: the creation of the cooled supersonic beam, interaction zone and detection. A separate chapter is devoted to the detailed description of the different setups of coils used to create the magnetic fields necessary to guide and to decelerate the particles of the beam.The third part is devoted to the experimental results and their direct interpretation using the equations of motion in Stern Gerlach forces. Simulations are presented to embody the interpretations. We present results about the deceleration of metastable argon and neon atoms. These results validate the significance of the addition of a uniform magnetic field defining a global adiabatic quantization axis for all the particles in the beam. This realizes the decoupling between the precession of the magnetic moments and Stern Gerlach forces. The results demonstrate the polarization effect of the beam that depends on the direction of the added uniform magnetic field relative to the progressive wave of the magnetic field.Finally, the understanding and control of the dynamics of trapping at a given speed, acceleration and deceleration require decoupling between the transverse and longitudinal effects of the wave. These effects are clearly visible when the added uniform magnetic field limits the transverse effects of the progressive wave of magnetic field. The outlooks for the new Zeeman Stern Gerlach decelerator are numerous. A first result of trapping di-nitrogen metastable at 560m/s is presented and the road is open to decelerate paramagnetic molecules in pulsed supersonic jet. Deceleration free radicals and neutrons are also possible.