I find an ordinary "thank you" entirely inadequate to express my gratitude to my advisor, Prof. Guruswami Ravichandran, who knows the students better than themselves and has never been parsimonious in encouraging them to explore their scientific enthusiasm and giving them his full support. I am indebted to him for introducing me to the enchanting world of experimental mechanics and providing the maximum freedom to let me pursue my interest in metallicglasses. In my moments of straying and wandering along the way, I especially thank him for bringing me back to focus, and helping me to distinguish important issues and interesting problems with a discerning and candid manner. Over four and half years I learned from him the great qualities of being not only a good scientist but also a responsible and considerate person. I wish I could bring his integrity, generosity, and sharpness with me in my future path.
The calculations are performed via the finite element method using the ABAQUS TM pack- age. A sequentially coupled thermal-stress analysis is conducted which involves the heat transfer calculations first. These yield the temperature evolution which is used in the sub- sequent mechanical analysis. Here, an implicit assumption is present, namely that the heat transfer and mechanical problems are uncoupled. This assumption is justified by the severe cooling during quenching which overwhelms heat generation in the sample due to inelastic deformation. The finite element mesh consists of a line of axis-symmetric elements along the radius of the composite. The interface between the two phases is assumed intact at all times. The elements of both phases are biased toward the interface. The kinematic boundary condition, in addition to axis-symmetry, is zero axial deformation at the bottom surface of the mesh. To impose the long cylinder condition, the nodes at the top are con- strained to move together in the axial direction (generalized plane strain). The analysis also includes the removal of the stainless steel elements after cooling to room temperature for the first two cases above (case 1: steel is elastic, and case 2: steel is elastic-perfectly plastic). This way, a direct comparison with the residual stresses measured in the BMG core can be performed.
It has been reported that crystallization proceeds either via an increase in temperature, which is induced by localized deformations, or a change in the chemical short-range order (CSRO) in the glassy state around the shear bands in the limited condition of a very large permanent strain. 3,10,11) However, little is known about whether crystallization occurs dynamically during a usual inhomogeneous deformation, i.e. the propagation of shear bands in the BMGs with high GFA. Deformation-induced nanocrystallization during quasi-static nanoindentation has been recently reported for Zr-Cu-Ni-Al- Ti BMG. 12) However, a change in the excess free volume
strength and high elastic limit.[99–103] However, most bulk MGs are brittle and fail catastrophically under tensile load because there are no effective plasticity mechanisms to suppress the sudden initiation and propagation of shear bands or cracks.[102–104] This lack of plastic deformability of MGs significantly limits their use in structural applications. Conventional toughening methods, for instance, making a composite with metallic phases, could improve the ductility of MGs, but the use of ductile and weak second phases usually degrades the yield strength.[105– 107] Due to the recent interests in nanotechnology, the nanoscale size effects on the mechanical properties of MG materials have been extensively studied,[108– 110] and it has been reported that nanosized MGs can exhibit extensive ductility without sacrificing the yield strength even under tensile stress.[111–113] Uniaxial mechanical tests of Pd-Si, Zr-based, and Ni-P MG nanopillars revealed a brittle- to-ductile transition by size reduction down to ∼100 nm in diameter.[111–113] Surprisingly, Zr-based and Ni-P MG nanopillars exhibit necking and strain hardening in uniaxial tensile tests.[112–114] In addition, radiation damage experiments at the nanoscale shows that MG nanopillars do not suffer from embrittlement when subjected to radiation damage with a Ga+ ion beam.[113–116] This is in contrast to high-strength conventional metals, which become brittle when exposed to high- energy radiation.[117–120] Unfortunately, these emergent phenomena of excellent resistance to brittleness in MGs are available only at the nanoscale. In order to harness this “smaller is more ductile” behavior and to proliferate it onto materials with large-scale dimensions, it is necessary to create a macroscopic meta-material that consists of nanosized components. 2
Wire Composites: Based on previous studies of Vitreloy 1 / W fiber composites, Mo, Ta, and stainless steel (Fe) fibers were used in a Vitreloy 106 matrix to enhance its ductility. In-situ loading experiments with neutron diffraction were performed to monitor strain evolution of the reinforcements vs. applied stress, and these results were combined with Finite Element Modeling (FEM) to deduce the deformation behavior of the glass matrix. It was found that the reinforcements yielded first and started transferring load to the matrix which remained elastic throughout the whole experiment. Despite the significant increase in ductility with these fibers, those fiber composites were not suggested as good candidates in real applications due to their low yield strength. The unwanted annealing during the melt infiltration process seems to further lower the yield point of the fibers, and hence the composite. Most fibers also exhibited a weak interface with the matrix which reduced the load transfer between the two phases. Due to its high yield strength, high stiffness, and good interface strength, W fibers appear to be the best reinforcements for BMGs. Critical issues to be addressed in future fiber composite research include: (i) Determination of optimum fiber volume fraction; (ii) measurement of fiber-matrix interface strength for various systems to quantify the effect of the interface on overall composite performance.
temperature. It is well known that a high strain rate promotes the inhomogeneous plastic deformation in BMGs, which is characterized by the formation of strongly localized shear bands, followed by catastrophic fracture. A relatively low strain rate is employed in the present tests for the detailed study of the plastic deformation process in the two BMGs. The Be free BMG exhibits a elastic modulus of 81.3 GPa and elastic limit of about 2.0% followed by a distinct plastic ﬂow
of the four BMGs ranges from 0.64 to 0.73. Plastic deformation behavior of the BMGs at various loading rates was studied by nanoindentation. The results showed that the loading rate dependency of serrated ﬂow, which is related to the nucleation and propagation of shear bands, depends strongly on the homologous temperature. The alloys with relatively high homologous temperature exhibit an increase in ﬂow serration with increasing loading rate, whereas, the alloys with low homologous temperature exhibit prominent serrations at low rates. No distinct shear band is observed around the indents for all alloys after nanoindentation at all the studied loading rates. Alternately, shear band pattern are characterized through macro-indentation, which shows that shear band spacing decreases with the increase of the homologous temperature.
during casting, (2) variations in residual stress, and (3) composition fluctuations in a given alloy. Each fracture specimen was produced via vacuum injection casting into copper mold from separately prepared ingots. Injection temperature during the casting process was not strictly controlled so the cooling rate could vary between specimens. In turn, this will leads to differences in the configurational state of the sample and associated free volume distribution . Residual stress develops during the casting process due to the hightemperature gradients which arise during sample cooling and solidification. Residual stress is known to affect fracture toughness significantly . According to Aydiner et al. , an 8.25mm thick Vitreloy 1 plate cast in a copper mold exhibited -25 to -30MPa surface compression and +10 to +13 interior tension. Their model suggested significant residual stress decreases with decreasing casting thickness. The casting thickness used in this study is 2.5mm. Aydiner et al. also showed that the compressive surface stresses were confined to a relatively thin surface layer. To reduce residual stress effects in the present work, ~10% of the surface layer was removed by grinding. This should
Residual stress is known to affect fracture toughness significantly . Residual stress develops during the casting process due to the hightemperature gradients which arise during sample cooling and solidification. Compressive stress develops in the surface while tensile stress develops in the interior. The development of the plastic zone could be influenced by the residual stress. Generally, heat treatment below glass transition temperature is used to anneal out residual stress. In this study, the surfaces of all the as-cast test specimens were ground off in an attempt to remove the compressive region of the residual stress and produce a certain degree of relaxation. According to the viscoelastic model of Aydiner et al. , an 8.25mm thick Vitreloy 1 plate is estimated to develop up to -230MPa surface compression and +90MPa interior tension. However, an actual experiment by these authors revealed that a copper mold cast piece with the same thickness exhibited only -25 to -30MPa surface compression and +10 to +13 interior tension. In addition, the model suggested significant residual stress decreases with decreasing casting thickness. The casting thickness used in this study is 2.5mm. Their results also indicate that the compressive surface stresses are confined to a relatively thin surface layer. Therefore, removing ∼10% of the surface layer is believed to reduce the residual stress to an insignificant level. Indeed,
Recently, a number of multi-component alloy systems, which makes the formation of glass phase possible at relatively low cooling rates (1–100 K/s), were discovered due to their higher thermal stability against crystallization. 1–3) Bulkmetallicglasses (BMGs) have been developed for structural applications utilizing their high strength, large elastic deformation limit, and superior corrosion and wear resistance at room temperature. Although the bulk of a cm- order dimension can be produced, the size to be used for engineering and structural application ﬁelds is still not enough. Therefore, the eﬀorts to relieve the size limit and to improve the workability of BMGs are needed to fabricate the components made of BMGs. 1,2)
Although, a vast majority of bulkmetallicglasses break shortly after yielding at ambient temperature, the deformation behavior of bulkmetallic glassy alloys is under intense investigations at present owing to their high strength. At relatively low homologous temperatures the inhomogeneous plastic flow of glassy alloys occurs by propagation of shear bands  which are 10–20 nm thick and make steps on the surface up to several micrometers in height . A strongly localized shear deformation at room temperature [65,66] limits practical application of such materials since a shear event may trigger a crack nucleation and rapid fracture. However, as illustrated in Figure 2, considerable apparent ductility was observed in several specific bulkmetallic glassy alloys [67,68].
A series of Cu-Hf-Ti alloys prepared by rapid solidiﬁcation of the melt and by copper mould casting were studied in the present work. Alloy ingots were prepared by arc-melting mixtures of pure metals in an argon atmosphere. An indication of the cooling rate obtained was determined using an Al-4.5 wt%Cu alloy. Cooling rates varied from 540 K/s for the centre section of a 4 mm die to 885 K/s for the outside wall section of the 2 mm die. The glass-forming ability, structure and thermal stability of Cu-Hf-Ti glassy alloys were studied by X-ray diffraction (XRD), differential scanning calorimetry (DSC) and differential thermal analysis (DTA). Bulk glass formation was observed for the Cu 64 Hf 36 , Cu 55 Hf 25 Ti 20 and Cu 56 Hf 25 Ti 19 alloys, with
1.1 The Urge for a New Technology VI-2 1.2 Problems of Current Lead-Free Solder Technology VI-3 1.2.1 Higher Temperature Requirement VI-3 1.2.2 Wetting, Spreading, and Bonding VI-4 1.2.3 Thermal, Mechanical, and Electrical reliability VI-4 2. BulkMetallic Glass as Solder Joints VI-6
Amorphous alloys suﬀered from three major issues, i.e., a limitation of product size and lacks of workability and weldability. Recently, amorphous alloys having a high glass- forming ability and a wide supercooled liquid region have been attracted increasing attention because of the high interests on basic science and engineering applications. The new amorphous alloys are called bulkmetallicglasses (BMGs). 1) They have superior mechanical properties such as high strength at ambient temperature and high-strain-rate superplasticity above the glass transition temperature. 2) Large BMGs can be fabricated by casting of the melt and by consolidation of the glass powders. 3,4) These results demonstrate that the BMGs have solved two major subjects of a limitation of product size and a lack of workability. However, the problem of welding has been left unsolved. In 2001, it has been reported that Pd 40 Ni 40 P 20 BMGs were
high-frequency ultrasound vibrations at elevated temperature, and found that the structural anomaly is induced by ultrasound vibrations in a supercooled liquid region. The electromagnetic acoustic resonance and resonant ultrasound spectroscopy methods were employed to measure the resonant spectra and ultrasonic attenuation coeﬃcients. When the glassy samples are subjected to sub/low-MHz ultrasound vibrations during heating process, the crystallization is accelerated around their glass transition temperatures, and with this abrupt structural change, irregular - shaped internal friction peaks appear. From the standpoint of ultrasonic echography, the glass transition and crystallization temperatures are considerably lowered by ultrasound vibrations in the present measurements.
group as well as in nonferrous alloy groups such as Ln-, 1) Mg-, 2) Zr-, 3) Ca- 4) and Cu-based 5) systems. Fe-based BMGs are classiﬁed into the following four groups of Fe-(Al, Ga)-P- C-B, 6) Fe-(Zr, Hf, Nb, Ta)-B, 7) Fe-Co-Ln-B 8) and Fe-B-Si- Nb. 9) It has been pointed out that the BMGs with high glass- forming ability (GFA) satisfy the following three empirical component rules, 10,11) i.e., (1) multicomponent systems
alloy (10.7), we can realize the formation of the icosahedral local structure in Zr-Al-Ni ternary glassy alloy. Therefore, it is concluded that the icosahedral local structure stabilizes the supercooled liquid state and the elemental combinations of Zr + Cu and Zr + Al + Ni are necessary for the icosahedral local structure formation. We also suggest that bulk glass- forming ability is attributed to the diﬀerence of stability of the icosahedral local structure with the combination of constitutional elements. The satisfaction of three component rules for high GFA leads to the high stability of icosahedral local structure by retarding the atomic rearrangements for the formation of stable crystalline phase.
BulkMetallicGlasses (BMGs) have been drawing increasing attention in recent years due to their scientific and engineering significance. A great deal of effort in this area has been devoted to developing BMGs in different alloy systems. BMGs based on certain late transition metals (e.g., Fe, Co, Ni, Cu) have many potential advantages over those based on early transition metals. These advantages include even higher strength and elastic modulii, and lower materials cost, to name but a few that are highly preferable for a broad application of BMGs as engineering materials. Nevertheless, these ordinary-late- transition-metal-based BMGs generally have quite limited glass-forming ability (GFA). In particular, for the Ni-based and Cu-based alloys reported prior to this research, the maximum casting thickness allowed to retain their amorphous structures is only ~2 mm (or lower) and ~5 mm (or lower), respectively.
Unlike in the inhomogeneous shear deformation, the generation and distribution of the Turnbull-Cohen free volumes are homogeneous as shown in Fig. 6. The relative free volume change reaches 0.4% at about 2% mean strain, slightly smaller than that predicted from elasticity using the Poisson ratio (0:34), and 1.5% when the overall elongation reaches 10% where the necking starts to form. In addition, we observed that the free volumes are distributed randomly in the sample; and most importantly, these seemingly randomly distributed free volume regions actually develop preferred orientations along the maximum resolved shear stress. As the deformation continues, both size and amplitude of the free volume regions increase. The local shear regions form in tandem with the formation of these free volume regions. At a critical (mean) free volume, necking occurs. Note that the critical free volume in tension deformation is smaller, about one-half of that at the peak stress in the shear deformation.
in the solute composition range less than 10 at%, and can be ignored because of a widely accepted concept that the nearly 10 at% solute atom is required for the formation of metallicglasses in a binary system. Thus, we reached the conclusion so far that the larger value of the exponent, at nearly n ≈ 50, is appropriate for evaluating ∆ H cavity for for binary systems.