Fabrication and Machining

of

Bulk Metallic Glass

for

Airborne Gravity Gradiometry

By

Kevin Mark Cole

A thesis submitted in conformity with the requirements

for the degree of Masters of Applied Science

Graduate Department of Materials Science and Engineering

University of Toronto

© Copyright by Kevin Mark Cole (2016)

ii

Fabrication and Machining of Bulk Metallic Glass for

Airborne Gravity Gradiometry

Masters of Applied Science (2016)

Kevin Mark Cole

Graduate Department of Materials Science and Engineering

University of Toronto

Abstract

Bulk metallic glass is an intriguing material ideally suited for use as a flexure in an airborne

gravity gradiometry. Successful fabrication of Zr56Ni20Al15Cu5Nb4 was achieved using arc

melting and suction casting. The effect of oxygen and microalloying Nb into this alloy

composition was investigated. It was determined that oxygen in solute form is much more

detrimental than as an oxide with respect to glass forming ability. Through microalloying Nb, a

high glass forming region was observed between 2 – 4 at.% Nb. Studies on crystallization

kinetics revealed that upon heating these amorphous samples, a multi-step phase transformation

pathway can be observed. Lastly, electrochemical micromachining (ECMM) and abrasive water

jet machining (AWJM) were shown to be effective techniques which can be used to shape

BMGs after casting without inducing crystallization. ECMM parameters were investigated to

optimize the micron-machining process. AWJM demonstrated that fast cutting could be achieved

with smooth surface finishes and good dimensional tolerance.

iii

Acknowledgements

I would like to sincerely thank and acknowledge the supervision, support, and guidance of

Professors Steven Thorpe, Donald Kirk, and Chandra Veer Singh throughout this thesis. They

provided me with the encouragement and knowledge required to fulfill my master’s degree.

They have been invaluable and taught me how to be a better all-around researcher moving

forward. I am truly thankful for their help.

I would like to acknowledge and show my appreciation towards Professor Keryn Lian, Haoran

(George) Wu, Sal Boccia, Abdolkarim Danaei, and the Department of Materials Science and

Engineering. I am extremely grateful for all their time, aid, and expertise as this work would not

have been possible without them.

I would like to thank Yuri Savguira, Qing (Alex) Ni, Josh Igel, Sina Sedighi, Jonathan Leung,

and Mohit Kumar for their assistance and mentorship.

Finally, I would like to thank my family and friends for their constant support and

encouragement – you are what keep me going.

iv

Table of Contents

Abstract ........................................................................................................................................... ii

Acknowledgements ........................................................................................................................ iii

List of Tables ................................................................................................................................ vii

List of Figures ................................................................................................................................ ix

List of Appendices ....................................................................................................................... xiii

List of Acronyms ......................................................................................................................... xiv

List of Symbols ............................................................................................................................. xv

1.0 Introduction ............................................................................................................................... 1

1.1 Motivation ............................................................................................................................. 1

1.2 Thesis Objectives .................................................................................................................. 3

2.0 Literature Review: Bulk Metallic Glasses ................................................................................ 5

2.1 Glass Forming Ability ........................................................................................................... 6

2.2 Processing of Bulk Metallic Glass ...................................................................................... 11

2.3 Crystallization Behaviour ................................................................................................... 12

2.3.1 Transformation Temperatures ...................................................................................... 12

2.4.2 Activation Energy ........................................................................................................ 13

2.4.3 Thermal Stability ......................................................................................................... 15

2.4.4 Structure ....................................................................................................................... 16

2.4.5 Minor Alloying ............................................................................................................ 17

2.5 Mechanical Behaviour ........................................................................................................ 20

2.5.1 Strength and Ductility .................................................................................................. 20

2.6 Corrosion Behaviour ........................................................................................................... 22

2.6.1 Zirconium-based BMG Corrosion Properties .............................................................. 22

2.6.2 Compositional Influences ............................................................................................ 23

2.7 Machining Bulk Metallic Glass .......................................................................................... 24

2.7.1 Electrochemical Micromachining ................................................................................ 27

2.7.2 Abrasive Water Jet Machining ..................................................................................... 29

3.0 Experimental Techniques ........................................................................................................ 31

v

3.1 Fabrication of Bulk Metallic Glass ..................................................................................... 31

3.2 Surface Characterization ..................................................................................................... 33

3.2.1 Optical Imaging ........................................................................................................... 33

3.2.2 Scanning Electron Microscopy .................................................................................... 33

3.2.3 Profilometry ................................................................................................................. 34

3.3 Differential Scanning Calorimetry ...................................................................................... 35

3.4 Composition Analysis ......................................................................................................... 36

3.4.1 Energy-dispersive X-ray Spectroscopy........................................................................ 36

3.4.2 X-ray Diffraction ......................................................................................................... 36

3.4.3 Inert-gas Fusion ........................................................................................................... 37

3.5 Transmission Electron Microscopy .................................................................................... 37

3.6 Electrochemical Techniques ............................................................................................... 39

3.6.1 Anodic Polarization ..................................................................................................... 39

3.6.2 Chronoamperometry .................................................................................................... 40

3.6.3 Electrochemical Micromachining ................................................................................ 41

3.7 Abrasive Water Jet Machining ............................................................................................ 45

4.0 Result and Discussion ............................................................................................................. 46

4.1 Production of Bulk Metallic Glass ...................................................................................... 46

4.1.1 Role of Niobium .......................................................................................................... 48

4.1.2 Effect of Oxygen .......................................................................................................... 50

4.2 Crystallization Behaviour and Thermal Stability ............................................................... 52

4.2.1 Enthalpy of Crystallization .......................................................................................... 53

4.2.2 Thermal Stability ......................................................................................................... 56

4.3 Electrochemical Characterization ....................................................................................... 62

4.3.1 Anodic Polarization ..................................................................................................... 62

4.3.2 Chronoamperometry .................................................................................................... 64

4.4 Electrochemical Micromachining ....................................................................................... 66

4.5 Abrasive Water Jet Machining ............................................................................................ 71

5.0 Conclusions ............................................................................................................................. 76

6.0 Future Work ............................................................................................................................ 78

References ..................................................................................................................................... 79

vi

Appendix A: Detailed Casting Protocol for BMGs ...................................................................... 88

Appendix B: XRD Database References ...................................................................................... 93

Appendix C: SEM Micrographs for As-cast BMG Rods ............................................................. 94

Appendix D: Kissinger Plots ........................................................................................................ 95

Appendix E: Ozawa Plots ............................................................................................................. 98

Appendix F: Anodic Polarization ............................................................................................... 101

vii

List of Tables

Table 1: Material property criteria for AGG flexure [1]. ................................................................ 2

Table 2: Comparison of key material properties between titanium, niobium and BMGs [1]. ....... 4

Table 3: List of various BMG compositions with corresponding Trg, γ, and Rc. ............................ 7

Table 4: Atomic size mismatch for common elements found in Zr-based BMGs. ........................ 9

Table 5: Heat of mixing for common elements found in Zr-based BMGs. .................................. 10

Table 6: List of various BMG compositions with corresponding Tg, Tx, ΔTx, and critical casting

diameter obtained at a heating rate of 20 or 40 K/min. ................................................... 16

Table 7: List of various Zr-based BMG compositions with corresponding E, σy, and ε. ............. 21

Table 8: SQFD table listing various machining criteria and techniques. Each parameter is

weighted from 1 to 10, with 10 being the most important. Machining techniques are

assigned a number from -9 to 9 based on how well it meets the criteria. The sum of all

the products is then tallied at the bottom indicating how well each technique meets the

given criteria. .................................................................................................................. 26

Table 9: Typical surface roughness and dimensional tolerance for EDM, ECMM and AWJM .. 27

Table 10: ECMM parameters and levels used to test and access machining efficiency on Zr-

based BMGs. ................................................................................................................ 44

Table 11: L9 Taguchi method outcome for ECMM experiments. ................................................ 45

Table 12: List of AWJM parameters used for machining of BMGs............................................. 45

Table 13: Atomic mismatch and heat of mixing for major constituents in the Zr56Ni20Al15Nb4Cu5

alloy system. ................................................................................................................. 47

Table 14: Casting conditions and compositions for suction cast 3 mm diameter rods of Zr-Ni-Al-

Cu-Nb alloys. ................................................................................................................ 47

Table 15: Oxygen content (wt.%) for each sample measured using inert gas fusion. .................. 51

Table 16: Characteristic thermal properties of Tg, Tx, and ΔTx, at 40 K/min for 3 mm diameter

as-cast amorphous rods. ............................................................................................... 56

Table 17: Peak temperature (Tp) at 10K/min and activation energies (Ea) of the first

crystallization peak measured through Kissinger Analysis for 3 mm diameter as-cast

rods and comparison to literature. ................................................................................ 61

Table 18: List of aspect ratios, surface roughness, depth, RMS and comments for ECMM holes.

...................................................................................................................................... 68

viii

Table 19: XRD data for phases present during casting and heating of Zr-Ni-Al-Cu-Nb BMGs.

The angle position (2θ), d-value (d), intensity (i), and crystal plane (hkl) for the first

three most intense peaks of each crystalline phase are given. ICSD card numbers are

given for reference. ....................................................................................................... 93

ix

List of Figures

Figure 1: Airborne Gravity Gradiometer (AGG) based on orthogonal quadrupole responder

(OQR) design [1]. ........................................................................................................... 1

Figure 2: List of various BMG systems with the maximum casting diameter that has been

achieved (modified) [11]. ............................................................................................... 8

Figure 3: Changes in Tg, Tx and ΔTx for Zr65Al10Cu15-xNi10Mx (M=V, Nb, Cr, Mo) [33]. .......... 19

Figure 4: Change in area of glassy phase for Zr65Al10Cu15-xNi10Mx (M=V, Nb, Cr, Mo) [33]. ... 19

Figure 5: Pitting corrosion seen in Zr-based bulk metallic glass. Right side is a magnified version

of left showing pitting behaviour [66]. ......................................................................... 23

Figure 6: Polarization scans comparing Zr56Ni20Al15Cu5Nb4 vs. Zr56Ni25Al15Nb4 and

Zr60Ni25Al15 in, from left to right, 1 N H2SO4, 1 N HCl, and 3 wt.% NaCl at room

temperature and open air (modified) [63]. ................................................................... 24

Figure 7: Schematic showing the general setup for ECMM [72]. ................................................ 27

Figure 8: ECMM hole made in Zr59Ti3Cu20Al10Ni8 BMG with corresponding surface profile

measured using a FRT MikroProf optical profilometer. Electrolyte is 2.94 M NaNO3,

ton = 100 ns and pulse voltage of 5V at room temperature [73]. .................................. 28

Figure 9: ECMM holes for a Fe65.5Cr4Ga4P12B5.5 BMG made in 0.1 M H2SO4 with 0.1 M

Fe2(SO4)3 showing the effect of different pulse times during machining [74]. ........... 29

Figure 10: Surface images of Zr52.5Cu17.9Ni14.6Al10Ti5 BMG after AWJM using a feed rate of

1000 mm/min. Right image is an enlarged in portion of the left showing broadening

marks [75]. .................................................................................................................. 30

Figure 11: Arc melter used for suction casting of BMGs. ............................................................ 32

Figure 12: BMG rod produced using arc melting and suction casting. ........................................ 32

Figure 13: Zygo white light profilometer station. ......................................................................... 35

Figure 14: FIB sectioning and lift out of BMG for thinning and TEM sample preparation. ....... 38

Figure 15: FIB thinning process used to produce Zr-Ni-Al-Cu-Nb BMG TEM samples. ........... 39

Figure 16: Schematic E vs. t and I vs. t for electrochemical scratch technique [78]. ................... 40

Figure 17: Rotating electrode ECMM setup. ................................................................................ 41

Figure 18: BMG sample used for rotating ECMM. ...................................................................... 42

Figure 19: Schematic and image of 10 μm diameter platinum tip used for ECMM [79]. ............ 42

x

Figure 20: CAD drawings and dimensions for ECMM setup. a) tool holder for holding Pt-

counter electrode and electrolyte jet. b) ECMM cell body. c) ECMM cell base. ....... 43

Figure 21: ECMM final cell set up used for machining of BMG samples. .................................. 44

Figure 22: XRD patterns taken for all as-cast samples (Compositions for I-VI given in Table 13).

Indexed phases were formed during casting process. ................................................. 48

Figure 23: High resolution TEM image (left) and selected area diffraction pattern (right) for

Sample III.................................................................................................................... 49

Figure 24: Effect of Nb concentration on the crystallization behavior noted in XRD, showing a

high GFA range from 2 – 4 at.% Nb. Dark regions correspond to Zr3NiO1.17 and

lighter region surrounding is the amorphous matrix and CuZr2. ................................ 50

Figure 25: DSC comparison for all as-cast samples at a heating rate of 10K/min. ...................... 53

Figure 26: Deconvoluted crystallization peaks for Sample III at a heating rate of 2.5 K/min with

corresponding XRD scans to identify existing phases at each peak step. .................. 56

Figure 27: Kissinger plot for suction cast 3 mm diameter rods of Zr-Ni-Al-Cu-Nb alloys using

peak temperature of the first crystallization peak. Error range is shown in the top

right corner. (Compositions I-VI as shown in Table 13). ........................................... 58

Figure 28: Example of the DDSC curves for Sample I used to find Tf. ....................................... 59

Figure 29: Comparison of activation energy found through Kissinger, Ozawa and Cheng analysis

with increasing Nb content (Compositions I-VI given in Table 13). ......................... 60

Figure 30: Anodic polarization scan for Sample III in 2.94 M NaNO3 at room temperature and

open air with scan rate of 1 mV/s. Section I is the passive region, II is the O2 reaction,

III is a change is oxidation states and IV is the pitting region. ................................... 63

Figure 31: Anodic polarization comparison between amorphous and crystalline samples. ......... 64

Figure 32: Current response to 1.785 V base during chronoamperometry. .................................. 65

Figure 33: Current response to 2.235 V base during chronoamperometry. .................................. 65

Figure 34: SEM image of ECMM hole at 3 V 10µs on: 100 µs off, no electrolyte jet. ............... 66

Figure 35: SEM image of ECMM hole at 3V, no pulse, no electrolyte jet. ................................. 67

Figure 36: SEM micrographs for ECMM holes. Numbers correspond to experiment number in

Table 9. ....................................................................................................................... 69

Figure 37: ECMM lines produce using experiment 3 parameters at 0.22 µm/s. .......................... 71

Figure 38: Trial 1 top surface of AWJM cut with magnification of edge on right. ...................... 72

xi

Figure 39: XRD pattern for Zr-Ni-Al-Cu-Nb BMG after Trial 1 AWJM process. ...................... 72

Figure 40: SEM cross section images after Trial 1 AWJM. Bulk (left) shows poor dimensional

tolerance as noted by cut line. Right image is a magnified image of the circle region.

..................................................................................................................................... 73

Figure 41: Trial 2 AWJM top surface and along with magnification of surface features in centre.

..................................................................................................................................... 73

Figure 42: Trial 2 exit (left) and entry (right) SEM images of top surface. Arrows indicate

direction of water jet machining. ................................................................................ 74

Figure 43: Cross section image of Trial 2 cut showing very good dimensional tolerance across

surface. ........................................................................................................................ 74

Figure 44: Detailed schematic of the Compact Arc Melter MAM-1 arc melter. .......................... 91

Figure 45: Image of vacuum chamber assembly and gas control panel. ...................................... 91

Figure 46: Images of a) the copper base plate detached from the vacuum chamber and b) normal

view of the base plate. ................................................................................................. 92

Figure 47: Image of a) split copper mould and b) bottom piece. .................................................. 92

Figure 48: SEM micrographs for cross-sections for as-cast rods listed in Table 12 taken at 500x

magnification to compare crystalline features. Sample I, IV and VI show clear

dendritic crystal formation. Sample II, III and V do not show any indication of

crystals; however, XRD indicates oxides are present in Sample II that cannot be seen

in SEM. ....................................................................................................................... 94

Figure 49: Kissinger plot for Sample I (Table 13). ....................................................................... 95

Figure 50: Kissinger plot for Sample II (Table 13). ..................................................................... 95

Figure 51: Kissinger plot for Sample III (Table 13). .................................................................... 96

Figure 52: Kissinger plot for Sample IV (Table 13). .................................................................... 96

Figure 53: Kissinger plot for Sample V (Table 13). ..................................................................... 97

Figure 54: Kissinger plot for Sample VI (Table 13). .................................................................... 97

Figure 55: Ozawa plot for Sample I (Table 13). ........................................................................... 98

Figure 56: Ozawa plot for Sample II (Table 13). .......................................................................... 98

Figure 57: Ozawa plot for Sample III (Table 13). ........................................................................ 99

Figure 58: Ozawa plot for Sample IV (Table 13). ........................................................................ 99

Figure 59: Ozawa plot for Sample V (Table 13). ....................................................................... 100

xii

Figure 60: Ozawa plot for Sample VI (Table 13). ...................................................................... 100

Figure 61: Anodic polarization scans for amorphous Sample V at a scan speed of 1 mV/s in 2.94

M NaNO3 at room temperature and open air. ........................................................... 101

Figure 62: Anodic polarization scans for partially-crystalline Sample IV at a scan speed of 1

mV/s in 2.94M NaNO3 at room temperature and open air. ...................................... 102

xiii

List of Appendices

Appendix A: Detailed Casting Protocol for BMGs………………………………………………88

Appendix B: XRD Database References…………………………………………………………93

Appendix C: SEM Micrographs for As-Cast BMG Rods………………………………………..94

Appendix D: Kissinger Plots ……………………………………………………………………. 95

Appendix E: Ozawa Plots ……………………………………………………………………….. 98

Appendix F: Polarization Curves …………………….…………………………………………101

xiv

List of Acronyms

AGG: Airborne Gravity Gradiometry

AWJM: Abrasive Water Jet Machining

BMG: Bulk Metallic Glass

BSE: Backscattered Electron

DDSC: Differential DSC

DSC: Differential Scanning Calorimetry

DTA: Differential Thermal Analysis

ECMM: Electrochemical Micromachining

EDM: Electric Discharge Machining

EDX: Energy Dispersive X-ray Spectroscopy

GFA: Glass Forming Ability

MG: Metallic Glass

OCP: Open Cell Potential

OQR: Orthogonal Quadrupole Responder

SAD: Selected Area Diffraction

SE: Secondary Electron

SEM: Scanning Electron Microscopy

TEM: Transmission Electron Microscopy

TTT: Time Temperature Transformation

XRD: X-ray Diffraction

xv

List of Symbols

A = constant (1/s)

D = effective diffusivity (# nucleus/cm3)

E = Young’s modulus (GPa)

Ea = activation energy (J/mol)

Eb = breakdown potential (V)

Ecorr = corrosion potential (V)

Erepass = repassivation potential (V)

G = shear modulus (GPa)

ΔG = Gibbs free energy (J/mol)

ΔG* = activation energy which must be overcome for formation of stable nuclei (J/mol)

ΔH = enthalpy of crystallization (J/g)

Hf = enthalpy of fusion (J/mol)

I = nucleation rate (# nucleus/cm3/s)

I = current density (A/cm2)

K = bulk modulus (GPa)

k = Boltzmann’s constant (J/K)

k = temperature-sensitive factor (1/s)

ko = constant (1/s)

n = exponent that reflects the nucleation rate and/or growth mechanism

R = gas constant (J/K/mol)

Rc = critical cooling rate (K/s)

Sf = entropy of fusion (J/mol/K)

T = temperature (K or °C)

xvi

Tf = inflection point of the first crystallization peak during DDSC (K)

Tg = glass transition temperature (K)

Tm = melting temperature (K)

Tp = peak temperature of first crystallization peak (K)

Trg = reduced glass transition temperature (unitless)

Tx = crystallization temperature (K)

ΔTx = supercooled liquid region (K)

V = voltage (V)

x(t) = volume fraction of crystalline phase (%)

β = heating rate (K/s)

γ = gamma parameter (unitless)

ν = frequency factor (unitless)

εy = elastic elongation (%)

σy = yield strength (MPa)

1

1.0 Introduction

1.1 Motivation

Gedex Inc. has produced a state of the art airborne gravity gradiometer (AGG) which resolves

gravitational gradients from land masses in order to explore and map out the location of various

elements and minerals below the Earth’s surface. The resolution is reliant on the pivot in the

orthogonal quadrupole responder (OQR) and is responsible for the spatial accuracy and

sensitivity of the signals. A schematic of the setup can be viewed in Figure 1 below with the

pivot being between Bar A and Bar B. The technique works by measuring the displacement of

the orthogonal beams around the pivot in response to gravitational gradients. When motion from

the aircraft is sensed the beams will move in the same direction and when a gravitational gradient

is detected they will move in opposite directions. Measuring these two different displacements

will allow for accurate measurement of where different masses exist subsurface. Since the OQR

is very sensitive to rotational and linear accelerations, it is important that the pivot is positioned

properly and has appropriate material properties to not cause extra movement.

Figure 1: Airborne Gravity Gradiometer (AGG) based on orthogonal quadrupole responder

(OQR) design [1].

2

Current AGG designs use either silica, niobium or titanium at cryogenic temperatures in order to

be able to distinguish and predict what minerals are present. None of these materials meet the

desired material requirements at room temperature and only at cryogenic temperatures can they

produce reliable signals. The desired material properties can be viewed in Table 1 below. These

material properties are important for achieving good signal to noise ratio and reproducibility

during operation. Having a high Young’s modulus will make the material stiff and is important

for lowering the natural frequency and maintaining geometric stability. The high tensile strength

and elastic deflection prevent the pivot from deforming due to the high stress and strains that are

applied. Additionally, having a high elastic limit will allow for the pivot to deflect more giving

higher signal to noise ratios during operation. The loss coefficient relates to the hysteresis of the

material when a stress is applied and removed. Therefore, it is important to have a low loss

coefficient such that the pivot will return to as close as possible its original position when no

gravity field is being sensed. Lastly, since the beam that will be attached to the pivot is made out

of Nb, it is important that the linear coefficient of thermal expansion of the pivot material be as

similar as possible to eliminate any thermal stresses or strains.

Table 1: Material property criteria for AGG flexure [1].

Material Property Initial Criteria

Young’s Modulus (GPa) ≥ 90

Ultimate Tensile Strength (MPa) ≥ 1800

Maximum Elastic Deflection (%) ≥ 1.8

Linear Coefficient of Thermal Expansion (m/m/ºC) Similar to Nb

Loss Coefficient / Internal Friction (Q-1

) 10-6

The drawback to the current design is that it requires a cryogenic chamber to be implemented

and carried in the aircraft along with the AGG. Additionally, most minerals and elements tend to

reside in remote areas making it very difficult to ship large quantities of liquid helium and

nitrogen and add extra weight to the aircraft, which increases fuel consumption and reduces

range/flight time. As such, the use of bulk metallic glasses (BMGs), also known as amorphous

metals, have been investigated as a substitute material for the pivot in the OQR due to their

3

enhanced mechanical and chemical properties at room temperature. Various BMGs have shown

similar mechanical properties at room temperature to niobium and titanium at cryogenic

temperatures. Therefore, the use of BMGs could not only eliminate the use of a cryogenic

chamber but also provide superior performance and resolution.

1.2 Thesis Objectives

Bulk metallic glasses have superior material properties over both titanium and niobium at room

temperature for the pivot in the OQR and similar to titanium and niobium when they are cooled

to cryogenic temperatures, as seen in Table 2. Therefore it is easy to see why there is a strong

push for BMGs, in particular zirconium-base BMGs, to be implemented due to their enhanced

material properties without the need of cryogenic temperatures. Even though BMGs are clearly a

superior material, they have yet to replace the existing titanium and niobium designs due to the

difficulty of both producing and machining BMGs. The absence of crystallinity is what gives

BMGs their increased material properties, but also what makes them difficult to machine once

cast. Most shaping techniques involve some form of mechanical deformation and/or elevated

temperatures, both of which can induce crystallization. Therefore, the objectives of the thesis are

to:

1. Select and reliably produce a BMG alloy with the mechanical and chemical properties at

room temperature desired for use in AGG

2. Investigate and understand the crystallization behavior of the produced BMG alloy for its

use in AGG

3. Determine and investigate appropriate machining techniques which do not lead to

crystallization of produced BMGs

4

Table 2: Comparison of key material properties between titanium, niobium and BMGs [1].

Material Property Titanium at

77K

Niobium at

20K

BMGs at

RT

Young’s Modulus (GPa) 115 125 146

Elastic Limit (MPa) 637 1150 2420

Ultimate Tensile Strength (MPa) 810 1175 2500

Maximum Elastic Deflection (%) 0.55 1.2 2.7

Loss Coefficient / Internal

Friction (Q-1

) 1E-06 9E-09 1E-07

Linear Coefficient of Thermal

Expansion (m/m/°C) 3.2E-06 2.0E-07 8.0E-06

Fatigue Limit (MPa) 54 at RT 75 at RT 1100

*BMG properties are taking from literature based on optimum values

5

2.0 Literature Review: Bulk Metallic Glasses

Traditional metals and alloys are usually formed and cast at relatively slow cooling rates which

allow atoms to rearrange during solidification and form periodic long range order that form into

crystals. The transition and rearrangement of atoms into crystals often causes defects to be

introduced, such as vacancies and dislocations, and often have negative effects on many of the

materials properties.

Metallic glasses (MGs) are a new class of materials that lack long-range crystalline order due to

fast cooling (rates ~106

K/s), which does not give sufficient time for atoms to rearrange in

crystals. Such high cooling rates are difficult to achieve in practice. Hence, newer forms of MGs,

known as bulk metallic glasses (BMGs), have become popular. They are at least a few

millimeters thick and require much lower cooling rates (<103K/s). The ability to produce BMGs

at certain cooling rate is referred to as the glass forming ability (GFA). The combination of high

GFA and cooling rates higher than the critical limit will freeze the random structural order along

with chemical short range order of the alloy in place and thus produce a noncrystalline alloy –

also referred to as amorphous metal or glassy alloy. Each alloy system has its own unique critical

cooling rate required for forming a fully amorphous alloy. Typically the cooling rate required to

produce BMGs is determined by the processing/casting technique and therefore allows alloys

with a lower critical cooling rate to form components with larger dimensions.

Zr-based BMGs are of particular interest due to unique combination of excellent mechanical and

chemical properties along with having a high GFA. Additionally, they are relatively low cost and

easy to fabricate in comparison to other forms of BMGs. Compared to their crystalline

counterparts, Zr-based BMGs typically show 2 – 3x higher strength and have high corrosion

resistance. They also have a toughness greater than most ceramics and elastic elongation similar

to some polymers. The following sections will detail the background knowledge behind

fabricating and machining Zr-based BMGs, and demonstrate why they are suitable materials for

use in AGG as a flexure material.

6

2.1 Glass Forming Ability

In order to first start understanding the glass forming ability (GFA) of an alloy system, it is

important to understand both the thermodynamic and kinetic aspects of crystallization [2]. In

1969, Turnbull [3] first suggested that a liquid with a glass transition temperature (Tg) over

liquidus temperature (Tl) equal to 2/3 becomes very difficult to crystalize and can only do so

within a very narrow temperature range. This lead to the determination that a high γ (Tx / (Tg +

Tl)), where Tx is the onset crystallization temperature, and a high reduced glass transition

temperature (Trg), which is the glass transition temperature over the liquidus temperature, are the

important indicators for determining the GFA of an alloy [4]. These indicators often tell how

easy or hard it is to form amorphous structure and is reflected in the critical cooling rate required.

Table 3 below lists the Trg, γ, and critical cooling rate (Rc) for various types of BMGs, namely

high glass formers. Au77.8Ge13.8Si8.4, a known poor glass former, is shown in Table 3 for

comparison and to demonstrate that metallic glass can only be produced for this alloy at

extremely high cooling rates. It can be seen that typically low values of Trg and γ require much

higher cooling rates in order to achieve a fully amorphous structure. It is also worth noting that

translating these parameters among different BMG systems are not very consistent. For instance,

Mg65Cu25Y10 and Pd78Cu6Si16 have the same γ, yet the Rc for the Pd-alloy is significantly higher.

Since Pd-based BMGs are naturally good glass formers, a value of 0.41 γ would indicate this

composition is not a good glass former as where this value in Mg-based systems would point

towards it being a moderately good glass former. Overall, it can be seen that Zr-based BMGs

very consistently show good GFA and require very low cooling rates that are comparable to even

much more expensive Pd-based BMG alloys.

7

Table 3: List of various BMG compositions with corresponding Trg, γ, and Rc.

Alloy Composition (at. %) Trg γ Rc (K/s) Reference

Au77.8Ge13.8Si8.4 0.48 0.33 3000000 [5]

Mg65Cu25Y10 0.55 0.41 100 [6]

Pd40Cu30Ni10P20 0.79 0.51 0.067 [7]

Pd78Cu6Si16 0.62 0.41 550 [8]

Zr65Al7.5Ni10Cu17.5 0.59 0.42 1.5 [9]

Zr41.2Ti13.8Cu12.5Ni10Be22.5 0.67 0.45 1.4 [10]

Measuring critical cooling rates is often a very tedious and difficult process. As such, the

maximum casting diameter is a more effective means to indicate the GFA of alloys. Figure 2

below shows the maximum rod diameter for a wide range of BMG alloys (Zr refers to

zirconium-based and could be any variation with zirconium being the primary element). This

maximum rod diameter is largely attributed to understanding the kinetics and thermodynamics of

the alloy as well as the casting technique used. Both of these will ultimately affect the

crystallization behaviour, physical properties, corrosion behaviour and mechanical behaviour of

the materials. Correlating Figure 2 to Table 3 shows that alloys which require lower critical

cooling rates will translate to greater diameters being achieved. Additionally, it shows that of all

the BMG alloys, Zr-based have the second highest achievable casting diameter only to the more

expensive Pd-based alloys.

8

Figure 2: List of various BMG systems with the maximum casting diameter that has been

achieved (modified) [11].

In addition to γ and Trg, there are additional factors which can influence the GFA. In fact, for

newer multicomponent BMGs, the intrinsic factors are known to play a more important role in

glass formation than external factors like cooling rate [8]. These intrinsic factors typically

include composition, cohesion among the metals, and the number, purities, and atomic size of the

alloying components [8]. Together they form what is known as the “confusion principle” [12],

which imply that increasing the number of alloying metals will destabilize competing crystalline

phases that would normally form upon cooling. This destabilization of the crystalline phase

allows for the melt to be more stable since the increase of alloying metals will lower the

tendency of the alloy to crystallize [8]. Each of these factors improves the GFA of a system by

limiting the amount of atomic movement during solidification such that the glassy phase stays in-

tact. Based on this, Inoue [13] suggested three empirical rules for glass formation in

multicomponent systems: (1) multicomponent systems shall consist of more than three elements,

(2) significant difference in atomic sizes with the size ratios above about 12% among the three

main constituent elements will promote higher GFA, and (3) negative heats of mixing among the

three main constituent elements [8] will improve GFA. These rules encapsulate the “confusion

principle” and correlate with both the kinetic and thermodynamic aspects for producing glassy

structures. Inoue’s criteria allowed quantifiable and tangible predictions of GFA. The effect of

each criterion is discussed in further detail below.

9

Increasing the number of elements in a system affects both the thermodynamic and kinetic

aspects of glass formation. From the thermodynamic perspective, having more elements will

increase the entropy of fusion (ΔSf) and increase the degree of dense random packing of atoms to

lower the enthalpy of fusion (ΔHf). Since Gibb’s free energy (ΔG) is the difference between ΔHf

and the product of temperature and ΔSf, the overall energy will decrease. This decrease in ΔG

will favour the formation of glassy phase over the competing crystalline phases. Additionally, it

will increase the solid-liquid interfacial energy which also favours the formation of glassy phase.

From the kinetic point of view, increasing the number of elements will increase the viscosity

which reduces both the nucleation and growth rates to supress crystal formation. This concept is

why Zr-based BMGs are often found in quaternary and quinary forms. Minor alloying is well

known to shift C-curves in time temperature transformation (TTT) to the right, allowing for more

time to cool and not produce nucleation. The selection of elements and quantities of minor

alloying elements are important factors for improving GFA [14, 15].

The second criterion is to have significant atomic size difference (above 12%) among main

constituents and is based on topological aspects such as structure and packing of atoms. Having

12% or more mismatch between major atoms will increase efficient packing of clusters and

increase the density of random packing of atoms within the supercooled liquid state. This will

then correlate with an increase in the liquid-solid interfacial energy and decrease atomic

diffusivity. Zr-based BMGs are often paired with varying amounts of Cu, Ni, and Al, among

other minor alloying elements. Table 4 below demonstrates that alloying these elements together

satisfies this rule and will lead to improved GFA.

Table 4: Atomic size mismatch for common elements found in Zr-based BMGs.

Element pairs Zr – Cu Zr – Ni Zr – Al

Atomic mismatch (%) 27 28 13

The third criterion, to have negative heat of mixing between major consistent elements, is

essential for mixing of atoms and formation of a homogenous glassy phase. Even though BMGs

do not have long-range structural order, they should still possess short-range chemical order. It is

10

important for BMGs to be homogenous such that the mechanical and chemical properties are

uniform throughout. Table 5 below shows the heat of mixing for Zr and other common elements.

Table 5: Heat of mixing for common elements found in Zr-based BMGs.

Element pairs Zr – Cu Zr – Ni Zr – Al

Heat of mixing (kJ/mol) -23 -49 -22

The capability to form a glass by cooling the melt from an equilibrium liquid is analogous to

suppressing crystallization within the supercooled liquid [8]. If we then assume steady-state

nucleation, the nucleation rate can be determined as a product of the thermodynamic factor

(driving force) and kinetic factor (diffusivity or viscosity) by the following equation:

𝐼 = 𝐴𝐷𝑒𝑥𝑝 [∆𝐺∗

𝑘𝑇] (1)

where: I = nucleation rate (# nucleus/cm3/s)

A = a constant (1/s)

D = effective diffusivity (# nucleus/cm3)

K = Boltzmann’s constant (J/K)

T = absolute temperature (K), and

ΔG* = activation energy which must be overcome for the

formation of stable nuclei (J/mol)

Thus by controlling these factors it is possible to achieve amorphous structures. For example,

with smaller values of ΔG the GFA will increase. For a larger degree of undercooling, the Gibbs

free-energy difference decreases as a result of stabilization of the undercooled melt due to an

increase in specific heat capacity from the lower free volume and short-range order in the alloy

melt [8]. Glass formers with lower critical cooling rates (deep eutectic) will have a lower Gibbs

free-energy than larger critical cooling rates with respect to the crystalline state [8]. Differential

thermal analysis (DTA) can indicate if an alloy is at a eutectic structure by a deep single phase in

the results which indicate small free volume and short-range order [8] [16].

11

Even with having all of the background knowledge on the thermodynamic and kinetic

requirements needed to form bulk metallic glasses, there still remains many challenges. Since it

is very hard to monitor and control every aspect of the casting process, it is possible that

materials which can oxidize do so and increase the chances of causing heterogeneous nucleation

[17]. The risk of oxidation can be reduced though fluxing techniques; however, it still creates a

non-ideal situation. Additionally, phase separation can also influence glass-forming ability and

introduce sites for nucleation [17]. While most of the obstacles that may arise during casting can

be avoided through various techniques and apparatus, a critical cooling rate is still required. As

such, the extraction of heat required to cool samples fast enough as the thickness increase for

bulk metallic glasses remains a key problem and influences various material properties

associated with the alloy.

2.2 Processing of Bulk Metallic Glass

Bulk metallic glasses can be processed through a variety of techniques, such as suction casting

into water cooled copper moulds, water quenching, and blow moulding, just to name a few [18].

The materials characteristics, such as formability and critical cooling rate, will determine which

process can yield a glassy structure. As such, the differences in the processing techniques could

potentially alter the material properties. In general, BMGs have high strength, large elastic strain

limit, high hardness, good soft magnetic properties, and excellent corrosion- resistance [13].

The most commonly used method for producing BMGs is to perform vacuum arc melting under

an argon atmosphere in conjunction with casting into water cooled copper moulds. This

technique is especially useful for casting Zr-based BMGs. Since oxygen can be very harmful to

the glass forming ability and casting process for Zr-based BMGs, it is important to melt under an

argon atmosphere [14]. Zr has a particularly high affinity for oxygen and thus the argon

atmosphere is required to reduce the amount of these impurities to achieve successful casting.

Even if a cooling rate well above the critical cooling rate is used, the presence of oxygen can

destabilize the supercooled liquid region and cause crystallization to occur [14].

Wang et al. [19] stated that monolithic BMGs show high compressive plasticity due to: (1) high

Poisson’s ratio; (2) chemical and microstructural inhomogeneity, including phase separation, soft

12

and hard regions with the same composition; (3) in situ nanocrystallization during deformation;

(4) more free volume by cooling faster or minor alloying. A high Poisson’s ratio seen in intrinsic

plastic BMGs will yield a low G/K (shear modulus over bulk modulus) and allow shear slip to

occur readily before the alloy fails by normal stresses [19]. However, this is counteractive for

achieving a high GFA. Additionally, since metallic glasses are amorphous and show plastic flow

instead of work-hardening, their tendency towards work-softening leads to shear localization

[17]. This localization of plastic flow into shear bands will limit the amount of plasticity it can

experience before failing catastrophically along a dominant shear band [17].

Therefore, the glass forming ability of the alloy will determine balancing improving plasticity

while still being able to cast as purely amorphous. In instances where the GFA cannot be

sacrificed, the process can be modified by using high pressure. Casting under high pressure tends

to suppress the nucleation and growth of the undercooled melts [8]. This suppression will

increase the alloys’ melting point and cause an increase in the undercooling of the liquid alloy

[8].

2.3 Crystallization Behaviour

Before discussing any of the material properties it is first important to understand the

crystallization behaviour of bulk metallic glasses. The improved material properties of BMGs

can be attributed to the elimination of crystals, thus understanding how they crystallize is

imperative for their use. While crystallization temperature (Tx) is not a material property itself,

due to it being sensitive to the heating rate used for measurement, it does provide a general guide

for the useful temperature range.

2.3.1 Transformation Temperatures

Differential scanning calorimetry (DSC) is a useful technique which can be used to not only

gather information on the material’s crystallization temperature (Tx), but also other important

transformation temperatures such as the glass transition temperature (Tg) and melting

temperature (Tm). The first change in slope seen in a DSC curve typically indicates the materials

glass transition temperature and signifies the change from a glassy state to a more relaxed state.

After the glass transition temperature, the material is now capable of forming crystals and a sharp

13

exothermic reaction will then occur each time there is a crystallization event. As such, the

number of exothermic peaks can then indicate how many crystallization processes the material

undergoes.

The Tg and Tx are important transition temperatures, especially when investigating crystallization

behaviour, but on their own do not provide information on the GFA. For example, the

Zr47(Cu4.5/5.5Ag1/5.5)46Al7 has a Tg of 702 K and Tx of 782 K, yet is not as good of a glass former

as Zr41.2Ti13.8Cu12.5Ni10Be22.5 which has a Tg of 625 K and Tx 705 K [4, 10]. However, a

parameter known as the supercooled liquid region can be obtained from these transition

temperatures and will be further discussed in Section 2.4.3. It is also important to note that

increases in heating rate will also shift these values higher and is important to compare data

obtained using the same heating rate.

2.4.2 Activation Energy

Studying the exothermic peaks observed during DSC also allows for the estimation of activation

energy. For BMGs, activation energy is most commonly calculated using either Kissinger’s

method [20] or Ozawa’s method [21]. Kissinger’s model assumes that with a change of heating

rate the temperature of maximum weight loss rate changes as well and that the change depends

solely on the activation energy for a given reaction [20]. For Ozawa’s method, it assumes that

when the DSC curve reaches its peak, the degree of reaction is at a constant and will be

independent of heating rate [21]. In order to derive the activation energy, both these methods

commonly use the peak temperature from the first crystallization event in DSC, but can also use

Tx. The Kissinger equation can be written as [20]:

𝑙𝑛𝑇𝑝2

𝛽=

𝐸𝑎𝑅𝑇

+ 𝑙𝑛𝐸𝑎𝑅

− ln 𝜈 (2)

where: Tp = peak temperature of the first crystallization peak (K)

β = heating rate (K/s)

Ea = activation energy (J/mol)

ν = frequency factor (s/s), and

R = gas constant (J/K/mol)

T = temperature (K)

14

The activation energy is then taken to be the product of the gas constant and the slope of the

linear fitting line when plotting ln (Tp2/β) versus 1/T. Alloys having a shallow slope, i.e. with a

high activation energy, will have greater thermal stability. Ozawa’s equation can be expressed as

follows [21]:

ln 𝛽 =−𝐸𝑎𝑇𝑝

+ 𝑐𝑜𝑛𝑠𝑡𝑎𝑛𝑡 (3)

where: β = heating rate (K/s)

Ea = activation energy (J/mol), and

Tp = peak temperature of the first crystallization peak (K)

The activation energy for this equation is obtained through the product of the gas constant and

the slope of the linear fitting line when plotting ln(β) versus 1/T. Both Kissinger and Ozawa

models are based on the assumption of first order reaction kinetics during continuous heating in

DSC. When the reaction does not follow first order reaction kinetics or nucleation rate and

growth mechanisms are desired, Johnson-Mehl-Avrami-Kolmogorov (JMAK) analysis can be

performed [22] [23]. In order to perform this analysis, isothermal DSC scans at different

temperatures are conducted and can be used to determine the volume fraction of crystalline phase

(x(t)) at time t according to:

𝑥(𝑡) = 1 − exp(−𝑘𝑡𝑛) (4)

𝑘 = 𝑘𝑜exp(−𝐸𝑎𝑅𝑇

) (5)

where: k = temperature-sensitive factor (1/s)

ko = constant (1/s), and

n = an exponent that reflects the nucleation rate and/or growth

mechanism

Ea = activation energy (J/mol)

T = temperature (K)

15

In order to obtain x(t), the area under the peak of the isothermal DSC at different times is divided

by the total area. Determining x(t) as such allows for Equation 4 to be written as:

ln[−𝑙𝑛(1 − 𝑥(𝑡))] = ln 𝑘 + nln(𝑡) (6)

Equation 6 can then be used to plot ln[-ln(1-x(t))] vs. ln(t) in order to obtain a straight line

whereby the slope is n [24].

2.4.3 Thermal Stability

As previously mentioned, transformation temperatures such as Tg and Tx can be determined

through DSC and provide input into the thermal stability of BMGs. In general, BMGs show good

thermal stability as noted by the large supercooled liquid region (ΔTx), which is defined by the

crystallization temperature minus glass transition temperature (Tx – Tg). It has been noted the

presence of a large supercooled liquid region will also aid to retard crystallization, which can be

used as another indicator of good GFA [25]. Referring back to Table 3 will show that larger ΔTx

are also associated with lower critical cooling rates. It has been seen that increased viscosity and

dense random packed structure can also be observed with large supercooled liquid regions [26].

Increased viscosity of the alloy will reduce any long range diffusion and make it difficult for any

long range diffusion to occur, which is required for crystallization. Therefore, it is possible to

attribute the increase in supercooled liquid region and thermal stability to a rise in viscosity and

dense random packing [26]. Table 6 below shows a comparison of Tg, Tx and ΔTx for various

BMG systems. The heating rate has less of an effect on the supercooled liquid region since Tg

and Tx generally translate similarly with increasing heating rate.

Comparing Table 6 with Table 3 shows that ΔTx is high for samples that have low critical

cooling rates. This indicates that ΔTx is a useful parameter for measuring GFA and is easily

obtained through DSC. Other GFA parameters such as Trg and γ require more sophisticated

equipment in order to measure the melting temperature, but are more accurate. As seen in Table

6, Pd40Cu30Ni10P20 has a lower ΔTx than Zr65Al7.5Ni10Cu17.5 even though it has a lower critical cooling rate

and has been seen to produce larger casting diameters. However, for Zr-based BMGs this parameter is

16

particularly sensitive and the wide range of values gives better insight into the GFA and crystallization

behaviour. This is seen by comparing two similar Zr-Cu-Ag-Al alloys in Table 6 that show large

differences in ΔTx which have also been reported to correspond to the maximum casting diameter

achievable [4]. It should be noted that ΔTx works best for indicating GFA within the same alloy system.

When comparing the Zr45(Cu4.5/5.5Ag1/5.5)48Al7 alloy to the Zr65Al7.5Ni10Cu17.5, it can be seen that the

Zr45(Cu4.5/5.5Ag1/5.5)48Al7 alloy has a lower ΔTx but has been shown to produce a greater casting

diameter. In general, ΔTx is not a very robust parameter for measuring GFA, but is still useful when

comparing Zr-based BMGs, especially within the same system. Additionally, it is an easily obtainable

parameter to can also be used to investigate the crystallization behaviour.

Table 6: List of various BMG compositions with corresponding Tg, Tx, ΔTx, and critical casting

diameter obtained at a heating rate of 20 or 40 K/min.

Composition T

(K)

Tx

(K)

ΔTx

(K)

Critical

Size (mm)

Heating rate

(K/min)

Reference

Zr43(Cu5/6Ag1/6)50Al7 738 770 32 20 20 [4]

Au55Cu25Si20 348 383 35 5 20 [27]

Pd78Cu6Si16 656 696 40 1-3 40 [28]

Ti40Zr25Ni2Cu13Be20 599 644 45 14 20 [29]

Ca60Al30Ag10 483 531 48 3 20 [30]

Zr45(Cu4.5/5.5Ag1/5.5)48Al7 711 785 74 > 20 20 [4]

Zr41.2Ti13.8Cu12.5Ni10Be22.5 625 705 80 14 20 [10]

Pd40Cu30Ni10P20 572 670 98 72 20 [31]

Zr65Al7.5Ni10Cu17.5 625 750 125 16 40 [32]

2.4.4 Structure

Pairing DSC with x-ray diffraction (XRD) and/or transmission electron microscopy (TEM) can

also give valuable insight into the crystallization behaviour that is seen. These techniques can

determine information such as what phases are forming for each individual event, phase

17

transformations, and even the growth rates of crystals. XRD is a relatively inexpensive tool

which allows for the identification of phases present at each crystallization event as well as the

phase transitions. However, this technique does have its limitations and drawbacks. The XRD

database is made up of patterns for crystalline powders with known compositions which lead to

sharp well defined peaks with no structural texture. If crystallization occurs during casting, it

often occurs in a non-stoichiometric manner which can have texture to it. This result can make it

hard to identify what phase is being formed, in addition to having an amorphous matrix in the

background. Additionally, very fine or small quantities of crystals may not be detected in XRD

and cannot be used to confirm the sample is completely amorphous.

On the other hand, TEM which is much more expensive and intensive in sample preparation,

excels in these areas when compared to XRD. In order to confirm a sample is truly amorphous or

not, selected area diffraction (SAD) patterns can be performed. Diffraction patterns are formed in

the reciprocal space and will show a broad diffused ring when amorphous and sharp defined

ring(s) (or dots) for crystalline samples. TEM can also be used to determine crystal compositions

and structures for non-stoichiometric crystals using energy dispersive x-ray spectroscopy (EDS).

Since EDS gives the amount of each chemical present it can quantitatively determine what each

crystal phase is.

2.4.5 Minor Alloying

The above section on crystallization behaviour discusses methods and techniques for measuring

and defining what a good glass former is. One of the most effective ways to increase the GFA of

a system is to alloy it with other minor elements. While minor alloying can also have an effect on

the materials properties (specifically strength and ductility) it is most beneficial when attempting

to suppress crystallization [14]. In order to effectively improve the GFA through minor alloying,

the element that is being added should still meet Inoue’s three criteria. Naturally, adding another

element will increase the entropy of the system. It also increases the degree of dense random

packing in the liquid state which decreases the enthalpy and solid/liquid interfacial energy.

Specifically, minor alloying plays a big role on increasing the supercooled liquid region, and as

mentioned above, will retard crystallization and improve thermal stability [25].

18

The magnitude of the effect of minor alloying varies largely from system to system. In the case

of the Zr65-xNi10Al10Cu`5Nbx alloy, the largest casting thickness was found at 2.5 at.% Nb [33];

however, the largest casting diameter was seen at 5 at.% Nb in a very similar Zr62-

xNi12.6Al10Cu15.4Nbx alloy [34]. These results show that minor alloying is not straightforward and

has to be optimized in order to improve GFA. This is largely due to presence and need of deep

eutectics which lowers the critical cooling rate required to form a fully amorphous alloy.

Furthermore, Figure 3 below shows the effect various alloying elements have on the Tg, Tx and

ΔTx within the Zr65Al10Cu15-xNi10Mx (M=V, Nb, Cr, Mo) alloy system. It can be seen that both

the amount and element have a great impact on Tg and Tx. The addition of Mo has little effect on

the Tx, but greatly increases the Tg and subsequently reduces the ΔTx. The opposite can be said

about V, until higher concentrations are added. The addition of Cr is seen to be negative on both

the Tg and Tx making it a poor minor alloying element for improving GFA. Lastly, Nb addition

shows that for low alloying amounts the ΔTx is not significantly affected, indicating a good

minor alloying element for improving GFA. These observations are further demonstrated in

Figure 4 where the amount of glassy phase was measured with increasing alloying elements. The

addition of V and Mo show signs of improved GFA to begin, but quickly decrease. The initial

rise in GFA is most likely a result of the higher entropy within the system, but as the

concentration is increased it will have a negative effect on the system. Addition of Cr was not

able to form any glassy structure which is to be expected by the results seen in Figure 3 for Cr

addition. For Nb addition, a broader range of concentrations was observed to enhance the GFA

with the most benefit being observed at 2.5 at. %. Even though ΔTx slowly decreased with

increasing Nb concentration, minor amounts are able to greatly improve the GFA of the system.

Overall, ΔTx is a good indicator of GFA, but does not always show improvements when minor

alloying. Minor alloying aids to increase GFA by the before mentioned intrinsic properties which

may not always be reflected in ΔTx.

19

Figure 3: Changes in Tg, Tx and ΔTx for Zr65Al10Cu15-xNi10Mx (M=V, Nb, Cr, Mo) [33].

Figure 4: Change in area of glassy phase for Zr65Al10Cu15-xNi10Mx (M=V, Nb, Cr, Mo) [33].

20

2.5 Mechanical Behaviour

The biggest opportunity with BMGs comes from their unique and superior combination of

mechanical properties. The strength of crystalline materials comes from the density of

dislocations, and is also what limits them from reaching their theoretical limit. In contrast, BMGs

do not have any crystalline features and are thus able to achieve much higher strengths. A high

elastic limit will allow for the flexure to avoid plastic deformation and give reproducible signals.

Large elastic elongations will improve the signal to noise ratio for sensing a wider range of

signals and differentiating them. Lastly, the absence of crystalline features also gives BMGs a

low loss coefficient which minimizes elastic hysteresis to improve repeatability.

2.5.1 Strength and Ductility

Most BMGs show compressive strengths of at least 1 GPa, while others have been seen to

exceed even 5 Gpa. Table 7 below summarizes a list of various Zr-based BMG alloys and the

range of key mechanical properties that can be observed during compression testing. Since the

flexure will only be operating within the linear elastic region, only properties corresponding to

this region are listed. These properties were determined through compression testing for similar

strain rates of 1 – 5 x10-4

s-1

.

It can be seen from Table 7 that the yield strength and Young’s modulus typically increase with

the complexity of the alloy. Adding elements, especially ones that are strong in their crystalline

form, aid to increase the overall strength. However, the increased strength can only be measured

in compression due to the brittle nature of BMGs. As shear bands begin to form in the plastic

region they are not inhibited by defects or grain boundaries allowing for rapid propagation once

they begin to nucleate.

Comparing results from Table 7 with Figure 3, it can be noted that Zr-based BMGs showed both

a high GFA and mechanical properties. The combination of these two factors is what makes Zr-

based BMGs very intriguing as they can be cast with large diameters and have exceptional

strength. In particular, Zr56Ni20Al15Cu5Nb4 meets the desired materials properties outlined in

Table 1 while having large casting diameters up to 25 mm.

21

Table 7: List of various Zr-based BMG compositions with corresponding E, σy, and ε.

Zr Based BMGs E (Gpa) σy (Mpa) ε (%) Reference

Zr60Cu18Al10Ni9Co3 97 1390 1.4 [35]

Zr55Al20Co20Cu5 92 1960 2.1 [36]

Zr65Cu15Al10Ni10 80 1570 2 [37]

Zr52.5Cu15Be12.5Al10Ni10 80 1750 2.2 [37]

Zr60Al20Ni20 78.2 1720 2.2 [38]

Zr60Cu20Al10Ni10 84 1790 2.2 [39]

Zr58Cu20Al10Ni10Ti2 85 1850 2.2 [39]

Zr50Cu40Al10 105 2000 1.9 [40]

Zr65Cu17.5Ni10Al7.5 92 1570 1.9 [41, 42]

Zr65Pd17.5Ni10Al7.5 86.1 1594 1.9 [41]

Zr55Cu20Al15Ni10 90 1850 2.1 [43, 44]

Zr52.5Cu20Al15Ni10Ti2.5 92 1800 2 [43, 44]

Zr58Cu22Al12Fe8 76 1710 2.25 [45]

Zr63.7Cu17.15Ni9.8Al7.35W2 91.5 1600 1.7 [46]

Zr53.9Cu29.4Al9.8Ni4.9W2 92.2 1800 2 [46]

(Zr55Cu30Al10Ni5)98.5Si1.5 87 1830 2 [47]

Vitreloy 1: Zr41.2Be22.5Ti13.8Cu12.5Ni10.0 96 1900 1.98 [48]

Zr48Be24Cu12Nb8Fe8 95.7 1600 1.7 [49]

Zr55Co25Al20-xNbx (x = 0, 2.5 and 5 at. %) 98 1960 2 [49]

Zr65Pd10Ni10Al7.5Cu7.5 85 1640 2 [50]

Zr60Cu15Pd10Al10Fe5 81 1701 2.1 [50]

Zr60Cu22Al10Au8 53 1740 3.3 [50]

Zr57Cu20Al10Ni8Ti5 68 1700 2.5 [51]

Zr52.5Cu17.9Ni14.6Al10Ti5 (Vit 105) 88.6 1760 2 [52, 53]

Zr60Cu15Al10Ni10Ti5 100 1752 1.8 [54]

Zr54.5Cu20Al10Ni8Ti7.5 110 1765 1.7 [55]

Zr63Cu14Al10Ni10Nb3 97 1949 2 [56]

Zr59Cu20Al10Ni8Ti3 78 1610 2.1 [57]

Zr65Cu12.5Ni10Al7.5Ag5 84 1650 1.95 [58]

Zr59Cu18Al10Ni8Ta5 85 1700 2 [59]

Zr37Be20.25Ti12.5Cu11.25Al10Ni9 119 1820 1.5 [60]

Zr35Cu17.9Hf17.5Ni14.6Al10Ti5 93 1900 2 [61]

Zr55Cu45 102 1750 1.7 [62]

Zr56Ni20Al15Cu5Nb4 100 1915 1.8 [63]

Ideal ≥ 90 ≥ 1800 ≥ 1.8 [1]

22

2.6 Corrosion Behaviour

While most of the interest with BMGs arises from their mechanical and magnetic properties,

BMGs also have superior corrosion behaviour compared to their crystalline counterparts. This

may be counterintuitive as BMGs have a metastable structure upon solidification that would be

expected to result in higher corrosion rates. The improved corrosion properties in BMGs can be

attributed to their chemical homogeneity, absence of crystal defects, and passivating films. Even

though BMGs possess no long-range structural order, they still maintain short-range chemical

order. The high cooling rates does not allow for solid-state diffusion and precipitates to form

which leads to chemical homogeneity. The absence of any precipitates or crystal defects, such as

grain boundaries, eliminates localized high-energy sites that could form a galvanic couple within

the alloy. Lastly, less passivating alloying elements is required for corrosion resistance since the

passive film will form uniformly on the surface.

2.6.1 Zirconium-based BMG Corrosion Properties

Due to the seemly endless combinations of BMGs, electrolytes, and concentrations; this section

will focus solely on the corrosion properties of Zr-based BMG due to its noted high glass

forming ability, good mechanical properties and the much lower cost compared to Pd-based

BMGs. Even though the corrosion performance of crystalline Zr is excellent in non-chloride

containing solutions, amorphous Zr-based BMGs perform even better. In chloride containing

solutions, such as NaCl, it was noted that the corrosion-resistance for BMGs was only slightly

better or equal to the crystalline form [64] [65]. However, the ability for amorphous Zr-based

BMGs to resist the onset of pitting was increased, indicating it is less susceptible to corrosion

[65]. In general, Zr-based BMGs act very similar to pure crystalline Zr. Upon polarization they

both show a strong passive film that is formed through anodization with a sharp breakdown

when pitting occurs. Even for Zr-based BMGs, pitting is still the dominant corrosion mechanism

which is found [66]. The main difference is that Zr-based BMGs are typically alloyed with

different elements which aid to passivate in various mediums and contribute to the resistance of

pitting in these alloys compared to the pure crystalline Zr form. Figure 5 below shows typical

pitting corrosion observed in Zr-based BMGs.

23

Figure 5: Pitting corrosion seen in Zr-based bulk metallic glass. Right side is a magnified

version of left showing pitting behaviour [66].

2.6.2 Compositional Influences

In order to improve the corrosion-resistance of Zr-based BMGs, elements such as Nb, Cr, and Ta

are added due to their known corrosion-resistance. When alloying these elements it is important

to take into account the effect it will also have on the glass forming ability and mechanical

properties. As such, when attempting to maintain other material properties and increase the

corrosion-resistance, Nb is the most commonly alloyed element since it is also used to enhance

GFA [14]. It has been observed that when replacing 20 at.% Zr with Nb in the Zr60-

xCu20Al10Ni10Nbx alloy, the corrosion rates in HCl were significantly reduced [67]. Additionally,

the inclusion of Nb in Zr-based BMGs not only improved the corrosion-resistance, but also the

prevention of pitting. This observation can be viewed in Figure 6 below as noted by the

increased Ecorr and breakdown potential (Eb). The addition of even small amounts of Nb causes

spontaneous passivation to occur and will aid in resisting pitting [67].

Typically Zr-based BMGs are paired with high Cu content (12.5 – 36 at.%) which has been

noted as the primary reason Zr-based BMGs are susceptible to chloride containing solutions and

H2SO4 [63]. A recent alloy (Zr56Ni20Al15Cu5Nb4) showed that reducing the overall amount of Cu

with Nb improved the corrosion-resistance, as well as improved the mechanical properties [63].

Figure 6 below compares polarization scans between this new alloy and traditional ternary

Zr60Ni25Al15 and a Nb containing quaternary Zr56Ni25Al15Nb4. It can be seen that in both H2SO4

24

and chloride-containing solutions, the low Cu alloy with Nb performs better than the other alloy

systems as noted by the increase in Ecorr and Eb. It is also worth noting that the passivating

current is typically lower when containing Nb. While the Ecorr and passivating current in NaCl is

better without the presence of Nb, there is no real passive behaviour and will pit very easily.

Since BMGs are chemically homogenous and do not contain defects, such as grain boundaries,

even a small amount of Nb will be able to prevent corrosion much more effectively than in the

crystalline form. The reduced corrosion resistant elements, along with the noted increase in Ecorr,

Eb and lower passivating currents all show Nb is an effective way to increase the corrosion

resistance of Zr-based BMGs.

Figure 6: Polarization scans comparing Zr56Ni20Al15Cu5Nb4 vs. Zr56Ni25Al15Nb4 and Zr60Ni25Al15

in, from left to right, 1 N H2SO4, 1 N HCl, and 3 wt.% NaCl at room temperature and open air

(modified) [63].

2.7 Machining Bulk Metallic Glass

As previously discussed, BMGs derive their superior mechanical and chemical properties due to

the absence of crystalline features. The presence of crystals, particularly dislocations and other

defects, lower the strength of the alloy and act as localized high-energy sites for corrosion. It was

also noted that the structure of BMGs are metastable and often begin to nucleate and crystallize

well below the glass transition temperature. Therefore, being able to find appropriate machining

techniques which do not alter the atomic structure is of key importance for utilizing BMGs in

industrial applications.

25

Even though the lack of atomic mobility upon casting minimizes the amount of shrinkage in

BMG, giving a very near net shape, there still may be a need for post-machining processes. In

fact, the lack of shrinkage often means any scratches or surface defects that are seen on the walls

of the mould will be transferred to the surface of the BMG. Additionally, some complex shapes

cannot be cast since the alloy needs to be molten right before casting and have the sufficient

cooling rate to form an amorphous structure. In both cases, machining may need to be conducted.

Micro-machining is preferred when attempting to remove surface defects. Macro-machining is

favoured when attempting to shape the final part. Regardless of the features that are to be

machined, the technique cannot use excess amounts of heat or deformation that would induce

crystallization.

Table 8 below shows an SQFD table (software quality function deployment) listing the desired

machining criteria and different techniques. This technique is used to determine what is the best

solution based on the customer’s needs. Each criteria is weighted from 1-10, with 10 being the

most important and 1 being the least. Based on how well each technique meets the criteria it is

given a score from -9 to 9. The sum of all products for each technique is then tallied to determine

the best overall solution. In general, it is very important that the technique does not crystallize

the sample and can achieve high dimensional tolerance (no dimensional changes across the

surface) and low surface roughness (< 2µm). Since only a few parts are required, the overall

complexity and time is not as big of an issue. By utilizing this method is can be seen that the top

three machining techniques, in order, are electric discharge machining (EDM), abrasive water jet

machining (AWJM), and electrochemical micro machining (ECMM). Table 9 shows physical

values for surface roughness and dimensional tolerance for these three techniques. Since electric

discharge machining was already being investigated [68], electrochemical micromachining and

abrasive water jet machining were chosen as potential techniques in this study. Research

conducted to test the validity of these techniques is discussed below.

26

Table 8: SQFD table listing various machining criteria and techniques. Each parameter is weighted from 1 to 10, with 10 being the

most important. Machining techniques are assigned a number from -9 to 9 based on how well it meets the criteria. The sum of all the

products is then tallied at the bottom indicating how well each technique meets the given criteria.

Machining

Criteria Weight

Milling/

drilling

Cutting

(spindle)

Laser micro-

machining EDM ECMM

Abrasive

water jet

Surface Finish 9 1 3 -9 -3 -1 9

Complexity 1 9 3 -3 1 -9 -1

Precision 8 -9 -3 9 1 3 -1

Crystallization 9 -3 -1 -9 1 9 3

Shaping Control 4 -9 -3 9 3 -1 1

Reproducibility 7 -3 -9 3 9 -1 1

Time 2 3 9 -1 1 -9 -3

Equipment/

Industry 5 1 3 -9 9 -1 -3

Total -127 -45 -83 113 53 89

Surface Finish – surface roughness (Ra)

Complexity – how difficult the machining process is to perform

Precision – the size and accuracy (tolerance)

Crystallization – likelihood of crystallizing during machining

Shaping control – freedom of motion during machining (i.e. 1D vs. 3D machining)

Reproducibility – ability to replicate features

Time – time to machine

Equipment/ Industry – availability and infrastructure of machining equipment and process

27

Table 9: Typical surface roughness and dimensional tolerance for EDM, ECMM and AWJM

Technique Surface Roughness (μm) Dimensional Tolerance (mm) Ref

EDM 0.4 – 0.8 ± 0.0025 – 0.13 [69]

ECMM 0.03 ± 0.005 [70]

AWJM 0.016 – 1.3 ± 0.03 – 0.05 [69, 71]

2.7.1 Electrochemical Micromachining

Electrochemical micromachining (ECMM) is a promising technique for machining localized

micron-size features. The technique works by pulsing either voltage or current between a tool

(such as a platinum electrode) and the sample (workpiece) in the solution to be machined, as

seen in Figure 7 below which shows a schematic set up of the system. ECMM is currently used

in industry for crystalline materials mainly because there is no tool wear, has high machining

rates and accuracy.

Figure 7: Schematic showing the general setup for ECMM [72].

Due to the success of ECMM with a wide range of crystalline materials, its feasibility as a

capable technique for patterning BMGs needs assessment. Since this technique uses neither

28

temperature nor deformation during the machining process there is no fear that it will crystallize

upon machining. The machining process is solely driven by selectively corroding fine areas to

remove material. As discussed in Section 2.5, BMGs are able to show better corrosion resistance

and resist pitting much better than crystalline forms. Paired with the ability to easily and quickly

repassivate have made it difficult to successively use ECMM. Work done by Koza et al. [73] has

shown that it is possible to machine Zr-based BMGs; however, the machining process is less

than desirable. Figure 8 shows a micrograph of a corresponding ECMM holes produced along

with the depth profile. The results demonstrate that the machining process for Zr-based BMGs is

irregular in both radius and depth. The irregular shape is most notably attributed to the

repassivating during pulse off times and corrosion products being formed [73]. More work on

finding optimal machining parameters and appropriate electrolytes need to be performed to

improve the machining process and achieve uniform holes similar to seen for various crystalline

materials.

Figure 8: ECMM hole made in Zr59Ti3Cu20Al10Ni8 BMG with corresponding surface profile

measured using a FRT MikroProf optical profilometer. Electrolyte is 2.94 M NaNO3, ton = 100

ns and pulse voltage of 5V at room temperature [73].

distance, µm

heig

ht,

µm

29

More recently, Horn et al. [74] was able to show that ECMM can be applied to Fe-based BMGs,

as seen in Figure 9 below. By using an aqueous electrolyte containing 0.1 M H2SO4 and 0.1 M

Fe2(SO4)3, along with ultra-short pulses, well defined holes with smooth surfaces were able to be

produced. Just as with Zr-based BMGs, the presence of a water based electrolyte and a

passivating surface, ECMM was difficult. However, with the addition of Fe2(SO4)3 in the

electrolyte, the machining process was improved. It was noted that the addition of the salt

inhibited the hydrogen evolution reaction and favored iron dissolution which is important for the

ECMM process [74].

Figure 9: ECMM holes for a Fe65.5Cr4Ga4P12B5.5 BMG made in 0.1 M H2SO4 with 0.1 M

Fe2(SO4)3 showing the effect of different pulse times during machining [74].

2.7.2 Abrasive Water Jet Machining

Another promising machining technique for BMGs is abrasive water jet machining (AWJM).

AWJM works by taking a mixture of water and abrasive which is shot through a pressurized

nozzle. This stream of water and abrasive creates a cutting jet that can be used to cut a wide

variety of materials including metals, polymers and ceramics. As such, this technique acts

similarly to wheel grinding during sample preparation, but with the ability to selectively machine

desired areas at high rates. This technique has been seen to be the more desirable method for

30

machining BMGs into complex parts without inducing crystallization [75]. In addition to this,

Wessels et al. [75] also noted that the surface temperature during AWJM only reached 63.1°C

which is well below any crystallization temperature.

While AWJM overall is a very intriguing machining technique, it still has some disadvantages.

Upon machining, broadening marks due to the expansion of the jet can be observed along the

surface (as seen in Figure 10) and reduces the overall dimensional tolerance across a cut sample.

This observation is even more prevalent in samples with round edges due to more tangents being

observed when the waterjet comes into contact and leaves the sample. To improve the surface

finish and dimensional tolerance, Loc et al. [71] performed abrasive water jet polishing. Through

varying different parameters, and using fine grit, they demonstrated that the surface roughness

could be reduced down to 0.016 µm [71].

Figure 10: Surface images of Zr52.5Cu17.9Ni14.6Al10Ti5 BMG after AWJM using a feed rate of 1000

mm/min. Right image is an enlarged in portion of the left showing broadening marks [75].

31

3.0 Experimental Techniques

The following sections outline the various techniques, characterization methods and machining

processes that were used over the course of this work. Details on sample preparation and

knowledge on the techniques are discussed.

3.1 Fabrication of Bulk Metallic Glass

Zr-based bulk metallic glass alloys with nominal compositions of (Zr57Ni20Al15Cu8)100-xNbx were

produced using arc melting and suction casting into a water cooled copper mould (Edmund

Büehler Compact Arc Melter MAM-1). Figure 11 shows the caster setup. The individual

elements, Zr – high purity (99.95 wt.% with 580 ppm O2), Zr – low purity (99.5 wt.% with 1400

ppm O2), Ni (99.98 wt.%), Al (99.999 wt.%), Cu (99.995 wt.%) and Nb (99.95 wt.%) were

obtained from Alfa Aesar and weighed using a Mettler AE260 before being placed in the arc

melter.

Before melting, the caster was evacuated down to a pressure of approximately 0.1 mbar using a

roughing pump and then filled back to atmospheric pressure using high purity argon (99.9999%).

This process also is used to create a pressure difference for suction casting between 1.01 bar and

0.1 mbar. This step was repeated five times for each sample to ensure the oxygen content

remained consistent. To ignite the plasma, the current is set to 6 on the arc melter (approx. 120

A) and then reduced to 3 (approx. 60 A). Each ingot was then melted four times for 15 seconds

under the high purity argon. Between each melting step, the ingot was flipped to achieve

improved chemical homogeneity before being finally suction cast into the mould after the fourth

melting cycle. The mould is 3 mm in diameter and 30 mm in length. This diameter complies with

all testing and sample preparation apparatuses required. The final product can be viewed in

Figure 12 below. A more detailed procedure can be viewed in Appendix A.

To study the effect of a pre-existing oxide on the crystallization behaviour, the casting process

was modified to a ‘cast-recast’ method. The cast-recast method involved taking a cast rod,

cutting it into several small pieces, and placing it back into the arc melter to be remelted for 15

seconds before being cast again.

32

Figure 11: Arc melter used for suction casting of BMGs.

Figure 12: BMG rod produced using arc melting and suction casting.

33

3.2 Surface Characterization

Sample preparation for surface techniques required sectioning the rod into smaller pieces

(roughly 3 mm in length) using a diamond blade cutter at 250 rpm and pressure <0.5 bar. After

sectioning, samples were followed by mounting in epoxy for grinding and polishing. Samples

were progressively ground on 400, 600, and 1200 grit SiC paper before being polished to a final

surface finish of 1 µm using diamond paste.

3.2.1 Optical Imaging

Optical imaging was used as a first means to access the surface for any surface defects or signs

of crystallization. While electron microscopes are typically needed to view any crystalline

features, in some extreme cases they can be observed using optical imaging. Imaging was also

used to keep a record of samples and be able to compare surfaces before and after various

experiments. Bulk imaging was done using a SZ-PT Olympus microscope while higher

magnification up to 2000x was done using a JENAPHOT 2000 microscope.

3.2.2 Scanning Electron Microscopy

For higher resolution imaging, scanning electron microscopy (SEM) was performed using a

Hitachi SU3500. SEM imaging was used for inspection and imaging of crystalline features and

machined features. Both secondary electron (SE) and backscattered electron (BSE) imaging was

performed. All SEM imagining was conducted under vacuum at 20 kV and beam current of 110

µA with a spot intensity ranging between 30 – 50. The working distance was set at 10 mm with a

tilt of 20° to ensure consistency between imagining and elemental analysis. All samples were

grounded to the stage using carbon tape to secure and eliminate any charging from the epoxy.

SE imaging is a technique whereby inelastic collisions from the incident beam interact with

valance electrons of the atoms. As such, this technique is very surface oriented (~100 nm) and is

used for exploring topographic conditions at the surface with high resolution. Specifically, it was

used for imaging and analyzing surfaces after machining. Not only is it important that the

machining technique does not crystallize the BMG sample, but also investigate the features it

creates. For ECMM, SE imaging can be used to assess the accuracy, size and shape of the holes

34

created as well as determine if pitting has occurred. The surface of AWJM is also investigated

using SE to access the flow of material, dimensional tolerance and roughness.

In contrast to SE, BSE imaging uses elastic collisions from the incident beam with nucleus of

atoms. Since this imaging is based off atomic nuclei it is able to show atomic contrast with

higher atomic numbers appears brighter. As such, this technique is used for investigating

crystallization even though the resolution of BSE imaging is lower than SE. Crystals that may

form during solidification will have a different atomic number than the amorphous matrix and be

differentiable through these elastic collisions. Additionally, since BSE can be generated from up

to 1 – 2 µm, it allows for better exploration of crystals.

3.2.3 Profilometry

Profilometry was performed to determine the surface roughness of samples and depths of holes

made in ECMM. Surface roughness and holes for ECMM samples were measured using a Zygo

white light profilometer, seen in Figure 13 below. This is a non-contact optical technique that

uses scattering from white light to determine surface profiles and roughness. The surface

measurements are gathered through coherence scanning interferometry which is a technique that

uses spectrally broadband visible light to achieve interference fringes. Through fringe

localization and interference fringe phase analysis, information about the surface topography is

obtained. This information was processed using MetroPro software to determine the depth of

holes as well as the average roughness (Ra) on the surface and inside the machined holes.

Samples were placed on the tilting surface and tilted until broad bands of light (fringes) could be

observed in the monitor. This was done to ensure proper scattering of the light and accurate

measurements could be obtained. An air pump was used to raise the surface and avoid any

vibrations during measurements. Using MetroPro software, the brightness was adjusted to be

between 90 – 100% before measurements. The area of analysis was typically around 0.5 mm by

0.5 mm with an objective lens of 20x magnification. Once the measurement is complete, a

roughness profile of the entire surface that was imaged will be shown in MetroPro. From this

plot, line measurements across the holes were used to determine the roughness and depth of each

hole.

35

For AWJM cuts, a mechanical profilometer was used (Mitutoyo Surftest SJ-410). This technique

used a stylus that is in contact with the surface of the sample and moved across the sample

measuring any vertical displacement. A skidless stylus with 0.75 mN force is used to maintain

contact with the surface throughout measurements. The stylus was a 2 µm diameter cylinder

which tapers off at 60° into a fine point. This stylus gives a resolution down to 0.000125 µm

with an 800 µm measuring range. Measurements are performed over the entire length (up to 25

mm for length) at a speed of 0.1 mm/s. Profilometry measurements were measured at least 3

times on each sample to obtain any error associated with measurements.

Figure 13: Zygo white light profilometer station.

3.3 Differential Scanning Calorimetry

Differential scanning calorimetry is able to measure the heat flow difference between a specimen

and a known reference pan. This heat flow gives insight into the thermal stability and kinetics

36

upon heating. The glass transition and crystallization temperatures of the rods were determined

through DSC (TA Instruments Q20) using heating rates (β) of 2.5 K/min, 5 K/min, 10 K/min, 20

K/min and 40 K/min under a high purity argon flow. Heating profiles used a single heating rate

from room temperature to 600° and allowed to air cool, unless otherwise stated. Enthalpy of

crystallization and activation energy (Ea) was measured with DSC as well. Each sample was

prepared using 15 – 20 mg of BMG and referenced against an empty hermetic aluminum pan. It

was noted that there was an error of ± 2.5°C associated with measurements.

3.4 Composition Analysis

A variety of compositional techniques; energy-dispersive x-ray spectroscopy (EDX), x-ray

diffraction (XRD) and inert gas fusion, were used to measure various elemental concentrations

throughout this work. These techniques were used to determine and/or quantify the bulk, crystal

and impurities present throughout the fabrication of BMGs.

3.4.1 Energy-dispersive X-ray Spectroscopy

Within SEM (Hitachi SU3500) energy-dispersive x-ray spectroscopy (EDX) was utilized to

determine the bulk composition of rods after casting. EDX does this by measuring the

characteristic x-rays that are produced as a result of the incident beam interacting with the

sample. Using 20 kV is sufficient for exciting x-rays in all measured elements. Additionally, a

beam current of 110 µA, working distance of 10 mm and tilt angle of 20° was applied. Since an

oxide layer forms during the casting process, the input composition varies slightly compared to

the output composition. As such, EDX can be used to determine the final composition and check

the rods for chemical homogeneity. However, due to the interaction volume that is created at 20

kV, compositions of small crystals that may have formed cannot be accurately determined due to

the relative size of the crystallite compared with the excitation volume (typically 2 – 5 μm).

3.4.2 X-ray Diffraction

To identify the crystalline phases x-ray diffraction (XRD) is utilized. XRD was performed using

a CuKα (0.15418 nm wavelength) which gives it much higher resolution (down to 2 nm) than x-

rays produced through an electron gun. Measuring the intensity of diffraction for various angles

37

produces a pattern which can be compared against known databases (such as PDF+4 and

inorganic chemistry structure database (ICSD) to identify which phases are present. XRD was

performed using a Rigaku Miniflex 600 with a CuKα to analyze the structure of the rods from 20-

80° at a scan speed of 6s per 0.03° step size and slit size between 2.5 – 5°. As a result of non-

equilibrium cooling used to produce BMGs, some phases formed upon casting are not

stoichiometric and cause identification difficulty. As mentioned, collected XRD patterns were

compared against databases which are recorded from stable phases and powders to eliminate

texture. Appendix B shows a list of references and ID numbers used for comparison. To identify

crystals, EDX in transmission electron microscopy would need to be performed. Lastly, pairing

XRD with DSC was used to identify phases forming during heating of amorphous structures.

3.4.3 Inert-gas Fusion

Inert gas fusion is a technique whereby a sample is heated to remove oxygen from within the

specimen. Once sufficient heat is given, the oxygen will be swept away by an inert carrier gas

such as argon or nitrogen, where it will then react with carbon to form CO2. The amount of

oxygen can then be determined by how much CO2 is detected. This technique was conducted to

determine the oxygen content in the as-cast rods (LECO ON736). As per industry standard,

preparation and testing was conducted according to ASTM E1409-13 [76]. To ensure

reproducibility, rods were sectioned into 0.14g samples using a diamond blade before testing.

Before each test, zirconium samples with known oxygen values were used to calibrate the

equipment and provide accurate readings of oxygen content. Through calibration, an error of ±

0.01 wt.% oxygen was obtained.

3.5 Transmission Electron Microscopy

Transmission electron microscopy (TEM) was performed to verify the amorphous structure.

While SEM and XRD and useful tools for imaging and identifying the presence of crystals, they

do have limitations. Very small amounts (< 1 wt.%) and size of crystals (typically below 2 nm),

can be missed by both of these techniques. In order to confirm the amorphous structure was truly

amorphous, selected area diffraction (SAD) and imaging in TEM was required. TEM operates

similar to SEM, but samples are thin enough to allow a high voltage electron beam to transmit

38

through the samples. In this work, a FEI Titan LB TEM microscope was used at 300 kV, which

gives an extraction voltage of 4300 V and gun lens of 3. The gun vacuum was set to 1 and

octagon vacuum below 20 in the operating software (Digital Micrograph). The first beam

condenser (C1) was set to 50 µm while the C2 was set to 100 µm, which are important for

controlling beam brightness. A magnification of 800x was used with a spot size of 3 – 4 for

imaging. This microscope was equipped with a CEOS image corrector to give a point resolution

of <0.1 nm. SAD patterns were produced at various points along the sample. Imaging was

performed using bright field and dark field detectors. Samples for TEM were prepared from bulk

specimens using focused ion beam (FIB) with a gallium beam at 5 kV. The bulk specimen was

placed inside the FIB machine where trenches are cut around the area of interest. This milling

process is done until the specimen is about 100 nm thick. In order to smoothen the surface and

obtain a flat surface, the sample is then tilted between 0.5-1°. This step is important since the

beam will broaden with depth and can cause thickness variations in the sample. Figure 14 below

shows the trenches formed (left image). The material on top is a tungsten film used to protect the

surface from any beam damage. Imaging was done using SEM mode in FIB. The sample is then

welded to a support such that it can be removed for further thinning as seen in Figure 14 (right

image). Once removed, FIB was then used to thin material until it becomes electron transparent.

Figure 15 shows the material being thinned in FIB. The sample is then placed on a carbon coated

Cu TEM grid such that it can be transported from the FIB to the TEM for imaging.

Figure 14: FIB sectioning and lift out of BMG for thinning and TEM sample preparation.

Tungsten

Film

39

Figure 15: FIB thinning process used to produce Zr-Ni-Al-Cu-Nb BMG TEM samples.

3.6 Electrochemical Techniques

BMG samples were analyzed before ECMM was conducted. All techniques were performed in

2.94 M NaNO3 solution at room temperature and open air. This concentration was chosen such

that the electrolyte would be slightly acidic and help dissolve any corrosion products that form.

Samples were progressively ground on 400, 600, and 1200 grit SiC paper. They were then

polished down to 1 µm using diamond paste, except for the ECMM samples. Before running

experiments, samples were left in solution for 30 minutes to reach open cell potential (OCP).

OCP measurements were conducted using an Ivium CompactStat potentiostat.

3.6.1 Anodic Polarization

Anodic polarization was performed to determine the electrochemical behaviour of Zr-based

BMGs in 2.94 M NaNO3 solution using a three electrode setup. Platinum mesh was used as the

counter electrode, the BMG was used as the working electrode and Ag/AgCl was used as the

reference electrode. Scans were produced by sweeping voltage at 1 mV/s between -0.9 to 3.5 V

vs. Ag/AgCl. Anodic polarization scans were also conducted using an EG&G Parc Model 616

rotating electrode set up. This set up was used to replicate and be able to compare with other

40

literature sources and help eliminate corrosion products [77]. Samples were rotated at 50 rpm

during polarization, while all other parameters were kept the same.

3.6.2 Chronoamperometry

Chronoamperometry was performed to do ‘electrochemical scratching’ on BMG rods to

determine the point at which they repassivate. Tests were performed using the same three

electrode setup mentioned in anodic polarization. The ‘electrochemical scratch’ technique is

performed by recording the current response when pulsing to a voltage where pitting occurs (i.e.

3 V), holding it for a short time (5 s), and then removing the voltage until it stabilizes. Each time

the voltage is pulsed, the biased voltage it returns to is set 50 mV higher than the last. The

repassivation potential (Erepass) is the voltage at which the corresponding current no longer

returns to zero. This value is often very similar to the breakdown potential seen during anodic

polarization, but is a more accurate means of determining what potential a protective layer will

no longer reform. A schematic of this process is shown in Figure 16 below.

Figure 16: Schematic E vs. t and I vs. t for electrochemical scratch technique [78].

41

3.6.3 Electrochemical Micromachining

First attempts to perform ECMM were done using a rotating electrode setup (see Figure 17

below). This setup was adapted and modified based on polarization studies for similar Zr-based

BMGs performed by Gebert et al. [77]. By rotating the sample at high rpm (500 – 2500), this

would allow for corrosion products to be removed upon machining and produce grooves along

the surface. Samples were prepared by placing as-cast rods into epoxy for grinding and to

provide suitable attachment onto the rotating head (seen in Figure 18). To perform micro

machining, the counter electrode used in this set up was a platinum tool 10 µm in diameter that

was encapsulated in glass. A schematic and image of this tool can be seen in Figure 19 below.

Figure 17: Rotating electrode ECMM setup.

42

Figure 18: BMG sample used for rotating ECMM.

Figure 19: Schematic and image of 10 μm diameter platinum tip used for ECMM [79].

Difficulties controlling the working distance and the inability to produce independent holes

limited the rotating electrode setup. Therefore, in order to produce holes to study machining

features and conduct electrochemical micromachining, a new electrochemical cell was designed,

produced and commissioned (see drawings in Figure 20 below and final product setup in Figure

21). The 10 µm platinum counter electrode is controlled by a motorized stage (Zaber Tech T-

LSM050A) in the x-y-z direction which allows for precise placement of the electrode tip within a

tolerance of 4 microns for micromachining. An outlet was attached such that electrolyte could be

recirculated and remove any corrosion products that may form and attach to the surface while

machining. The cell base is threaded and bolted to the stage to eliminate any motion of the cell

body during machining. Samples were leveled using a Newport tilt stage. This was done for the x

and y direction and was leveled to within 3 µm over the length of the sample. Once the sample

was leveled, the platinum counter electrode was place in the tool holder and brought to within 15

± 10 µm of the sample surface. After OCP was achieved, machining was performed for 30 – 60

minutes at a given spot. After one hole was performed, the tool tip was translated 2 mm away to

machine at a different location. This was done such that the electrochemical fields would not

interact with each other. For lines, the tip moved at a scan speed of 0.00022 mm/s.

43

Figure 20: CAD drawings and dimensions for ECMM setup. a) tool holder for holding Pt-

counter electrode and electrolyte jet. b) ECMM cell body. c) ECMM cell base.

44

Figure 21: ECMM final cell set up used for machining of BMG samples.

It was determined that the peak voltage, base voltage, duty cycle and flow rate were important

factors when micromachining. Table 9 below outlines these parameters and the setting used for

each parameter. To evaluate each of these parameters, the Taguchi method was applied and

outlined the experimental program that was used for testing. Table 10 shows the experimental

setup that was derived from the Taguchi method in order to test and access these various

parameters and levels.

Table 10: ECMM parameters and levels used to test and access machining efficiency on Zr-

based BMGs.

Control Factor Levels

1 2 3

A. Peak Voltage (V) 3 4 5

B. Duty Cycle (ton:toff) 1:1 1:5 1:10

C. Flow Rate (mL/min) 0.2 0.3 0.4

D. Base Voltage (V) 0 1.785 2.235

45

Table 11: L9 Taguchi method outcome for ECMM experiments.

Experiment Number A (V) B (Ton:Toff) C (L/min) D (V)

1 3 1:1 0.2 0

2 3 1:5 0.3 1.785

3 3 1:10 0.4 2.235

4 4 1:1 0.3 2.235

5 4 1:5 0.4 0

6 4 1:10 0.2 1.785

7 5 1:1 0.4 1.785

8 5 1:5 0.2 2.235

9 5 1:10 0.3 0

3.7 Abrasive Water Jet Machining

Two sets of trials were performed using AWJM. The first set of trials was performed using a

Flow Mach 3b waterjet. After clamping one side of the rod, the waterjet was passed though the

cross-section to make a cut. The second set of trials was using an OMAX Maxiem 0707 waterjet.

Instead of clamping the rod, for this trial the pieces were tapped into sacrificial blocks. The block

of aluminum was used to help reduce the defects that can arise when the waterjet comes into

contact and exits the rods, as noted in Section 2.7.2 and Figure 10. Parameters for both trials are

summarized in Table 11 below.

Table 12: List of AWJM parameters used for machining of BMGs.

Parameter Trial 1 Trial 2

Machine Flow Mach 3b Waterjet OMAX Maxiem 0707

Cutting speed 0.1” / minute

Working distance 0.1”

Abrasive 80 mesh garnet 120 HPX Barton hard rock

Nozzle 1 mm 0.5 mm

Water pressure 52 ksi

Fixture Clamp one side Tapped in aluminum

46

4.0 Result and Discussion

The results obtained through fabrication and machining of (Zr57Ni20Al15Cu8)100-xNbx bulk

metallic glass are presented. The production of bulk metallic glasses and the role of minor

elements/impurities is presented in Section 4.1. Section 4.2 shows an in-depth investigation into

the thermal stability and crystallization behaviour. The characterization of appropriate

electrochemical behaviour is shown in Section 4.3. Machining was performed on the micro and

macro scan. For micro machining, electrochemical micromachining results are presented in

Section 4.4. Abrasive water jet machining results are shown in Section 4.5 for macro machining.

4.1 Production of Bulk Metallic Glass

The material properties of BMGs are heavily dependent on the amorphous structure. As such, it

is important to understand the role of minor alloying elements on the GFA and crystallization

behaviour during the casting process. The alloy selected for fabrication is the newly discovered

Zr56Ni20Al15Nb4Cu5 by Li et al. [63]. This alloy was selected since it satisfies the required

mechanical and chemical properties required for use as a flexure material in AGG. Additionally,

referring back to Inoue’s criteria, this alloy is seen to meet the requirements and leads to the

enhanced GFA. Since it is a quinary system it meets the first criteria, and Table 12 below shows

the atomic mismatch and heat of mixing for major elements. It can be seen that all the atomic

mismatch are above 12% and heat of mixing are all negative except for Zr-Nb, which is a minor

alloying element. It has also been reported for similar Zr-Ni-Al-Cu-Nb systems, that too much

Nb will reduce the GFA and form quasicrystals [80]. Impurities, such as oxygen, can also have a

similar effect on the crystallization behaviour. Since Zr-based BMGs are particularly sensitive to

oxygen, it is also important to understand the role in which it plays during the casting process.

47

Table 13: Atomic mismatch and heat of mixing for major constituents in the Zr56Ni20Al15Nb4Cu5

alloy system.

Zr-Ni Zr-Al Zr-Cu Zr-Nb Ni-Al

Atomic mismatch (%) 28 13 27 13 14

Heat of mixing (kJ/mol) -49 -44 -23 +4 -22

Table 13 below lists the six 3 mm diameter as-cast (Zr57Ni20Al15Cu8)100-xNbx (from x = 1.9 to 4.3

at.%) alloy rods used to investigate the role of Nb and oxygen on the crystallization behaviour

during suction casting. The relevant casting parameters and resulting microstructure caused by

Nb and oxygen concentration are also listed. The ‘cast-recast’ method was previously explained

in section 3.1. Input compositions of Zr57Ni20Al15Cu8-xNbx were used. Due to an oxide layer

being formed on the top outer surface upon casting, the as-cast compositions were measured

using EDX (seen in Table 13). Additionally, having various crystallization pathways and

fabrication methods may also alter the final composition. Using EDX as a means of estimating

compositional variations of ± 0.5 at.% can be expected.

Table 14: Casting conditions and compositions for suction cast 3 mm diameter rods of Zr-Ni-Al-

Cu-Nb alloys.

Sample EDX Comp. (at.% ± 0.5%) Zr Purity Casting Observation

I Zr58.8Ni20.0Al14.5Cu4.8Nb1.9 High Suction Partially crystalline

II Zr55.4Ni21.3Al15.4Cu5.4Nb2.5 High Cast-recast Partially crystalline

III Zr55.5Ni20.6Al14.7Cu6.5Nb2.7 High Suction Amorphous

IV Zr55.0Ni20.3Al15.9Cu5.3Nb3.6 Low Suction Partially crystalline

V Zr54.9Ni19.9Al15.0Cu6.3Nb3.9 High Suction Amorphous

VI Zr54.5Ni20.7Al15.2Cu5.2Nb4.3 High Suction Partially crystalline

48

4.1.1 Role of Niobium

Through varying the Nb content, different crystallization behaviour was observed. Figure 22

shows the XRD patterns for all samples listed in Table 13. Analyzing the XRD patterns showed

that the GFA is not sufficient in Sample I where the Nb concentration was less than 2 at.%. The

resulting pattern shows the formation of Zr3NiO1.17 and CuZr2. Upon increasing the Nb

concentration to 2.7 at.% (Sample III), a broad pattern hump and no crystalline peaks is

observed, corresponding to an amorphous structure. The amorphous structure continues to be

formed up to 4 at.% Nb (Sample V). To confirm the amorphous structure seen in XRD and

corresponding SEM images in Appendix C, TEM imaging and diffraction has been performed

(seen in Figure 23 below). The left image is a higher resolution TEM image showing no long

range order and the typical ‘salt and pepper’ image indicative of an amorphous BMGs. This was

further confirmed through selected area diffraction pattern (Figure 23, right). The obtained

diffraction pattern showed broad diffused rings which further verified the amorphous structure.

These rings appeared this way since SAD patterns are viewed in the reciprocal space. Since there

is no long range order within amorphous samples and only various atomic clusters of different

orientations, the rings will appear broad and diffused.

Figure 22: XRD patterns taken for all as-cast samples (Compositions for I-VI given in Table 13).

Indexed phases were formed during casting process.

49

Figure 23: High resolution TEM image (left) and selected area diffraction pattern (right) for

Sample III.

Once the Nb concentration surpassed 4 at.% (Sample VI), signs of crystallization can be seen

again. Concentrations higher than 4 at.% Nb revealed the beneficial effects of Nb on GFA are

lost and once again Zr3NiO1.17 and CuZr2 crystals were observed, along with Nb0.94O0.06 crystals

due to the increased Nb concentration. These results suggest there is possibly a deep eutectic

point between 2 and 4 at.% Nb that helped to stabilize the liquid and prevent crystallization from

occurring, thereby creating an ideal casting range for this low Cu alloy system [81]. This ideal

casting range is consistent with results seen for (Zr0.65Al0.10Cu0.15Ni0.10)100-xMx by Inoue et al.

[33] and Zr56Ni20Al15Nb4Cu5 by Li et al. [63]. The high GFA region and crystallization pathways

are depicted in Figure 24 below which also shows Zr3NiO1.17 (dark crystals) and CuZr2 as the

lighter crystals in the amorphous matrix. SEM images for all six alloys can be viewed in

Appendix C.

50

Figure 24: Effect of Nb concentration on the crystallization behavior noted in XRD, showing a

high GFA range from 2 – 4 at.% Nb. Dark regions correspond to Zr3NiO1.17 and lighter region

surrounding is the amorphous matrix and CuZr2.

4.1.2 Effect of Oxygen

Even if the alloy contained a favorable amount of Nb (2 – 4 at.%), the presence of oxygen has

been reported to hinder its GFA. Many reports have found the role of oxygen was to inhibit the

GFA of Zr-based BMGs, but the form, or state, in which the oxygen is present, can also have a

great effect on the GFA [82, 83, 84, 85]. Within an alloy, oxygen can exist as either a solute

element or part of an oxide. These different forms will have varying effects on the crystallization

pathway upon casting. Therefore, it is important to understand how each will affect the casting

process and the severity on the GFA. Oxygen existing as a solute material and part of an oxide

was studied individually to determine their influence on the crystallization mechanisms.

In order to study oxygen as a solute material, the concentration was varied through the oxygen

content in the raw materials. High purity Zr (99.95 wt.% with 580 ppm O2), and low purity Zr

Zr3NiO1.17

CuZr2

Zr3NiO1.17

CuZr2

51

(99.5 wt.% with 1400 ppm O2) was used to produce samples with varying oxygen solute

concentrations. Two samples (IV and V) of similar compositions were cast using Zr with high

and low concentrations of oxygen. The slight variations in compositions measured from EDX

can be attributed to the crystallization pathways during casting. Oxygen values in the as-cast rods

were measured using inert gas fusion and are summarized in Table 14 below. These

measurements confirmed that the use of low purity Zr increased the oxygen content by two fold.

Higher oxygen concentrations lead to the crystallization of the supercooled liquid and formation

of metastable phases seen in Figure 22 and Appendix C. DSC was also used to confirm the

presence of these metastable phases (Figure 25, Sample IV). In DSC, the amorphous samples go

through multiple phase transformations as a result of not having pre-existing crystals. Therefore,

since Sample IV only shows one peak during DSC it further supports the existence of pre-

existing crystals. The presence of these metastable crystals destabilized the supercooled liquid

and reduced the GFA of the alloy (6). These findings are consistent with observations reported in

similar alloy compositions [82, 83, 84].

Table 15: Oxygen content (wt.%) for each sample measured using inert gas fusion.

Sample I II III IV V VI

Oxygen (wt.% ± 0.01) 0.06 0.08 0.08 0.13 0.07 0.11

To study the effects of a pre-existing oxide, the cast-recast method was utilized. Through casting

samples of similar composition (Sample II and III) with high purity oxygen and using the cast-

recast method, a solid oxide layer was introduced into the alloy. To confirm the oxide was

present within the rod, XRD was conducted as seen in Figure 22. The sharp peak at 29° is

attributed to the oxide layer (monoclinic ZrO2) that was formed during the first cast and

incorporated into the alloy during the re-cast. The pattern seen in XRD also shows signs of

crystallization, but is much less apparent than for Sample IV, the high solute content alloy.

Analyzing the surface with SEM was not able to detect any crystals which would further suggest

52

only slight crystallization occurred. Inert gas fusion results seen in Table 14 show the total

oxygen content between Sample II (no oxide) and Sample III (pre-existing oxide) show the two

are equal in value. This result confirms that the minor crystallization observed in XRD was due

to the pre-existing oxide and not a rise in solute oxygen. Therefore, it can be concluded that the

influence of oxygen on the crystallization behaviour is closely related to the form in which it is

present. Oxygen in the form of solute causes the supercooled liquid region to destabilize due to

the formation of metastable crystals, while oxygen as an oxide is relatively less harmful to

forming BMGs. Similar observations were previously reported [82, 83, 85].

It should be noted that the role in which oxygen plays in the fabrication of BMGs is more

complex than appears from the findings of this investigation. Even though current findings

suggest that the presence of oxygen in solute or oxide form reduces the GFA of the alloy, other

investigations have found that a certain amount of oxygen can be beneficial to the GFA [84, 85].

In these studies, it was shown that reducing oxygen below a certain level can cause

crystallization to occur and only when a certain amount of oxygen was present could a fully

amorphous structure be produced. For the results presented here, Sample I shows a high degree

of crystallinity even though it has the lowest oxygen content. Therefore, it can be concluded that

reducing the oxygen content in the alloy is not always the preferred strategy to improve GFA. In

such situations, minor alloying (Nb above 2 at.%) can be seen to be a more effective method to

stabilizing the supercooled liquid region and produce an amorphous structure.

4.2 Crystallization Behaviour and Thermal Stability

Even though the investigation of Zr-based BMGs was to ensure their thermal stability in an AGG

operating at room temperature, it is still important to determine the crystallization behaviour and

temperature range of thermal stability. Since the amorphous structure of BMGs is metastable,

and heat will be generated from flexing the material, it must have sufficient thermal stability to

not transform or crystallize to retain the same material properties throughout operation.

53

4.2.1 Enthalpy of Crystallization

To investigate the crystallization kinetics of these alloys, DSC scans were conducted with

varying heating rates ranging between 2.5 and 40 K/min. A comparison of DSC curves for the

alloys in Table 13 can be viewed in Figure 25 below at a heating rate of 10 K/min.

Figure 25: DSC comparison for all as-cast samples at a heating rate of 10K/min.

Figure 25 shows that amorphous samples (III and IV) have a clear glass transition temperature

followed by multi-step exothermic crystallization peaks. On the other hand, the glass transition is

not so obvious for as-cast partially crystalline samples (I, II, IV, and VI) and only a single

exothermic crystallization peak is observed. Amorphous samples showed a multi-step

crystallization mechanism which begins with the formation of cubic Zr2Ni (c-Zr2Ni) that

transforms into tetragonal Zr2Ni (t-Zr2Ni) and Zr5Ni4Al [86, 87]. Each of these phases produce a

separate peak as shown in the DSC curve. From each of the DSC curves, the enthalpy of

crystallization (ΔH) can be estimated and used to characterize the GFA and degree of

crystallinity in the alloys. Since partially crystalline samples already have crystallites present in

their microstructure which act as nucleation sites to lower the activation energy required for

54

further crystallization, the ΔH for these samples are typically lower than for fully amorphous

samples. However, comparing ΔH to amorphous samples can give insight to the degree of

crystallization that has occurred. Amorphous samples exhibit higher values of ΔH due to having

fewer nucleation sites and greater thermal stability. As such, amorphous samples will have

different crystallization pathways compared to partially-crystalline ones and therefore cannot be

used to directly determine the amount of crystalline features. Within amorphous samples, it has

been reported that greater enthalpy of crystallization also shows larger supercooled liquid regions

and improved glass forming ability [88]. Analyzing the enthalpy of crystallization for these

samples showed an average ΔH of 82 J/g for amorphous samples and 68 J/g for partially

crystalline samples. Both of these values are improvements over the enthalpy calculated for the

Zr55Cu30Al10Ni5 alloy by Lin et al [87].

Analyzing the enthalpy associated with each crystallization step in amorphous samples can give

a greater insight into the crystallization kinetics upon heating. As previously mentioned, each

exothermic peak seen in DSC can be attributed to a phase transformation. Therefore by

measuring the area under each peak it was possible to estimate the enthalpy of crystallization for

each phase and provide insight into the crystallization pathways and its overall contribution to

the enthalpy of the crystallization process. To determine the area under each peak, Gaussian

curves were fitted to each peak to deconvolute their individual contributions. An example is

shown in Figure 26 below. During increasing heating rates from 2.5 to 40 K/min, it was observed

that the first peak, corresponding to the formation of c-Zr2Ni crystals, only gradually increased

while the second peak (t-Zr2Ni) decreased and the third peak (Zr5Ni4Al) increase much more

than the first one. These trends indicate that the crystallization pathway will begin to favour peak

3 as heating rate is increased as well as demonstrate the importance of the first peak. The need

for the first peak indicates that c-Zr2Ni is required as a precursor before any further phase

transformations can occur. These results were supported through XRD analyses shown in Figure

26. To identify the phase present at each phase transformation, XRD was performed after DSC

was used to heat the sample up to the corresponding crystallization temperature (Tx1 for peak 1

and Tx2 for peak 2) and then air cooled down to room temperature. For peak 3, the sample was

taken from previous runs to 600°C to ensure complete phase transformation occurred. Figure 26

shows a sequential reduction in the presence of c-Zr2Ni as the sample is heated from peak 1 to 3.

55

The XRD spectrum for peak 3 shows t-Zr2Ni is still present even after heating to 600°C. These

results indicate there are two crystallization pathways c-Zr2Ni to t-Zr2Ni and c-Zr2Ni to

Zr5Ni4Al. As heating rate increases, the crystallization pathway will have sufficient energy to

favour the formation of Zr5Ni4Al and bypass the t-Zr2Ni phase. It is thus believed that these two

phases are the equilibrium phases that will form.

The present section examined the enthalpy of crystallization and correlated it to the

crystallization mechanisms of amorphous (Zr57Ni20Al15Cu8)100-xNbx alloys through the

combination of DSC and XRD techniques. Deconvolution of individual DSC peaks provided

new information on the role and stability of each phase transformation, while XRD was used to

identify the chemistry and structure of these phases. Sequential XRD patterns at each

crystallization peak also demonstrated the crystallization pathways that occur during heating.

It is also noted that upon heating, various forms of oxygen will form. Even though DSC was

performed under an argon atmosphere these phases still exist. This observation further

demonstrates just how sensitive this Zr-based BMG is to oxygen and its high affinity for it. Both

Nb and Zr are sensitive to oxygen; however, the Gibb’s free energy based on Ellingham

diagrams shown ZrOx is more favourable to form at these temperatures. Additionally, the higher

concentration of Zr in the alloy will lead to more oxide phases being formed with it compared to

Nb. It is also worth noting that t-ZrO2 typically doesn`t form until 1443 K, yet it is present even

in Peak 2 (around 775 K). This observation could be a result of having t-Zr2Ni present and

allows for the higher symmetry t-ZrO2 to form. As more time and temperature is given (Peak 3),

the expected m-ZrO2 then begins to form. The formation of this phase is lower in symmetry and

is the expected phase below 1443 K. Overall, these results further show that Zr-based BMGs are

extremely sensitive to oxygen.

56

Figure 26: Deconvoluted crystallization peaks for Sample III at a heating rate of 2.5 K/min with

corresponding XRD scans to identify existing phases at each peak step.

4.2.2 Thermal Stability

Materials which have good thermal stability are able to retard crystallization more easily.

Thermal properties such as supercooled liquid region and activation energy can be used to

indicate if a material has good thermal stability [89]. As previously mentioned, the supercooled

liquid region (ΔTx) is defined as the difference between the crystallization temperature (Tx) and

glass transition temperature (Tg) from DSC curves. Table 15 shows the Tg, Tx, and ΔTx for

amorphous samples (III and IV). It can be seen that the ΔTx for Sample III is slightly higher than

Sample V; however, this is still within the error range of DSC. As such, activation energy can

also be used to indicate the glass forming ability of each alloy.

Table 16: Characteristic thermal properties of Tg, Tx, and ΔTx, at 40 K/min for 3 mm diameter

as-cast amorphous rods.

Sample Composition (at.% ± 0.5%) Tg (K) Tx (K) ΔTx (K)

III Zr55.5Ni20.6Al14.7Cu6.5Nb2.7 722 771 49

V Zr54.9Ni19.9Al15.0Cu6.3Nb3.9 729 776 47

57

The activation energies were determined and explored using three different non-isothermal

equations; Kissinger, Ozawa and Cheng [20, 21, 90]. While Kissinger is the most common

method used to calculate activation energy, a comparison of activation energies helps to

determine accuracies and trends. Both Kissinger and Ozawa models assume first order reactions

which neglect nucleation and growth [91]. These models use the peak temperature of the first

crystallization event seen in DSC to derive the activation energy [90, 91]. Cheng`s method is

based on similar criteria but uses the temperature from the first inflection point whereby the

crystallization rate is at a maximum [90].

For Kissinger analysis, the activation energy is taken to be the product of the gas constant and

the slope of the linear fitting line when plotting ln (Tp2/β) versus 1000/T, as seen in Figure 27.

Kissinger’s model can be viewed in Section 2.1, equation 2. Alloys having a shallow slope, i.e.

with a high activation energy, will have a greater thermal stability and won’t be affected by

temperature as much. The error range associated with these calculations were derived by

performing tests over multiple samples at the same heating rate and is noted in the top right

corner of Figure 27. Once again, this range of variability was within that of the equipment,

indicating good reproducibility of crystallization behaviour for these samples. These results show

that samples with shallow slopes (III and V) have higher activation energies and will be more

thermally stable. Even though Sample IV shows signs of crystallization, it has a shallow slope

similar to Sample III and V. This result is most likely due to the degree of crystallinity within the

sample as it only shows one crystallization peak in DSC for all temperatures. Since Kissinger`s

model is based on the reaction rate at the peak temperature, this will translate very similarly to

amorphous samples with increasing heating rate. As such, the calculated activation energy using

Kissinger will give similar results to that of Sample III and V which are amorphous. Lastly,

linear regression showed each tested sample to have a R2 greater than 99%. Individual plots can

be viewed in Appendix B.

58

Figure 27: Kissinger plot for suction cast 3 mm diameter rods of Zr-Ni-Al-Cu-Nb alloys using

peak temperature of the first crystallization peak. Error range is shown in the top right corner.

(Compositions I-VI as shown in Table 13).

For the Ozawa method, the activation energy for this equation is obtained through the product of

the gas constant and the slope of the linear fitting line when plotting ln (β) versus 1000/T.

Ozawa’s method can be viewed in Section 2.1, equation 3. Individual plots can be viewed in

Appendix C. Lastly; Cheng’s equation is stated as [90]:

ln𝑇𝑓2

𝛽=

𝐸𝑎𝑅𝑇𝑓

+ 𝑙𝑛𝐸𝑎𝑅

− ln 𝜈 (7)

where: Tf = inflection point of the first crystallization peak (K)

β = heating rate (K/s)

Ea = activation energy (J/mol)

ν = frequency factor (s/s), and

R = gas constant (J/K/mol)

59

This equation is identical to Kissinger with the only difference being the use of Tf instead of Tp

in Equation7. To determine Tf value, differential DSC (DDSC) curves were produced. Figure 28

shows an example of DDSC curves for Sample I. Cheng’s equation was original meant for the

determination of activation in oxide glasses and not metallic glasses; however, Wang et al [91]

demonstrated its effectiveness for calculating activation energies for a Zr70Cu20Ni10 bulk metallic

glass. It should be noted that only fully amorphous samples were tested to reflect the structure of

oxide glasses.

Figure 28: Example of the DDSC curves for Sample I used to find Tf.

The results of estimating the activation energies for all six as-cast samples using Kissinger,

Ozawa and Cheng’s method are compared in Figure 29 below. Analyzing these results show that

Kissinger and Ozawa follow a similar trend when estimating activation energy for both

amorphous and partially crystalline samples. The activation energy from Ozawa’s equation are

consistently higher than Kissinger which is also noted in similar alloys [90, 91]. Activation

energies calculated for amorphous samples (III and V) using Cheng’s equation show equal or

lower values compared to Kissinger and Ozawa. This observation was also seen in the

Zr70Cu20Ni10 alloy by Wang et al. [91]. However, results were inconsistent when using Cheng’s

60

equation to calculate activation energy for samples with pre-existing crystals. Although Cheng’s

method gives lower activation energy values for amorphous samples, it is unable to capture the

effect of pre-existing crystalline phases on thermal stability. Refereeing back to XRD patterns in

Figure 22, Sample IV shows a high degree of crystallinity while Cheng’s equation estimates a

higher activation energy for this sample over the amorphous ones. This inability of Cheng’s

equation in this regard suggests that Kissinger and Ozawa analyses shall be considered as more

appropriate models for calculating activation energy for BMGs.

Figure 29: Comparison of activation energy found through Kissinger, Ozawa and Cheng

analysis with increasing Nb content (Compositions I-VI given in Table 13).

The similarity in Kissinger and Ozawa indicates that the Kissinger model indeed provides a good

estimation for the activation energy of BMGs, even in presence of crystals in small volume

fraction. As such, the Ea values for Samples I – VI are given in Table 16 below. It is seen that the

activation energies calculated through Kissinger for amorphous samples (III and V) are greater

than the (Zr0.645Ni15.5Al11.5Cu8.5)100-xNbx alloys by Iqbal and Akhter, and comparable to the

Zr60Cu20Al10Ni10 alloy by Zhuang et al. [92, 93]. These results suggest that Sample III and V are

more thermodynamically stable and therefore possess high GFA.

61

Table 17: Peak temperature (Tp) at 10K/min and activation energies (Ea) of the first

crystallization peak measured through Kissinger Analysis for 3 mm diameter as-cast rods and

comparison to literature.

Sample Composition (at.% ± 0.5%) Tp (K) Ea (kJ/mol)

I Zr58.8Ni20.0Al14.5Cu4.8Nb1.9 767 239

II Zr55.4Ni21.3Al15.4Cu5.4Nb2.5 789 250

III Zr55.5Ni20.6Al14.7Cu6.5Nb2.7 772 316

IV Zr55.0Ni20.3Al15.9Cu5.3Nb3.6 795 306

V Zr54.9Ni19.9Al15.0Cu6.3Nb3.9 775 307

VI Zr54.5Ni20.7Al15.2Cu5.2Nb4.3 791 270

[92] Zr0.645Ni15.5Al11.5Cu8.5 --- 260

[92] (Zr0.645Ni15.5Al11.5Cu8.5)98Nb2 --- 232

[93] Zr60Cu20Al10Ni10 --- 326

As stated at the beginning of Section 4.2, the supercooled liquid region and activation energy

could be correlated to the thermal stability and GFA for BMGs. Comparing the supercooled

liquid region between amorphous samples showed only minor changes that were within the error

range of DSC. However, after also comparing activation energies between these two systems it

can also be seen that Sample III has a higher activation energy than Sample V. Together these

results suggest that Sample III will have a slightly higher GFA [8, 17, 86]. Sample III contained

a Nb concentration of 2.7 at.% and is consistent with similar alloys for increasing GFA [33, 34].

This finding suggests there is a possibility that the critical diameter of the new alloy casted by Li

et al. [63] could be even further increased if cast at 2.7 at.% Nb. Further work is required to find

the optimal concentration within the high GFA region, but this result does demonstrate there is

room to further enhance this alloy system.

It should also be noted that comparing activation energy between amorphous samples (III and

IV) and crystalline samples (I, II, IV, and VI) shows the activation energy for the first

crystallization peak is higher than that for amorphous ones. The higher activation energy leads to

amorphous structures being more stable as they would require more energy to crystallize. This

result is most likely attributed to the lack of heterogeneous nucleation sites present in amorphous

samples. Pre-existing crystals can act as effective nucleation sites.

62

4.3 Electrochemical Characterization

It is important to understand the electrochemical behaviour of this new Zr-based BMG before

ECMM can be performed. It is important to know information such as the passivating behaviour,

breakdown potential and repassivation potential. The breakdown potential is of particular

importance since it is above this potential in which ECMM can be performed and maintain

constant machining rates.

4.3.1 Anodic Polarization

Performing anodic polarization gives a general snapshot into the electrochemical behaviour of a

sample in an electrolyte. Figure 30 below shows the polarization scan for Sample III in 2.94 M

NaNO3 solution at room temperature and open air. Current was measured in response to an

increasing voltage scan at 1 mV/s and then converted to current density (A/cm2) based on the

original active area. Analyzing these results shows this alloy follows the typical Zr-based BMG

trend in that it forms a stable passive layer followed by breakdown (Eb). The Ecorr (-128 mV vs.

Ag/AgCl) for this alloy system is seen to be higher than the high-Cu concentration

Zr57Cu15.4Al10Ni12.5Nb5 (-144 mV vs. Ag/AgCl), but still lower than Zr59Ti3Cu20Al10Ni8 (-54 mV

vs. Ag/AgCl) and Zr55Cu30Al10Ni5 (-44 mV vs. Ag/AgCl) [77]. As the potential continues to

increase after Ecorr the sample transitions into region I where spontaneous passivation occurred.

The presence of Nb and lack of defects makes it very easy to passivate as noted by the absence

of a Tafel slope on the anodic side and steady current densities throughout the region. Since the

electrolyte is water based, increasing the potential to region II showed the onset of the oxygen

evolution reaction to occur and lead to a “permeability” of the passive layer. This permeability is

what causes a rise in the current densities during region II. Increasing the potential to region III

will cause a change in oxidation states which allowed for the sample to form a new passive layer.

However, once Eb is exceeded (region IV) pitting occurred and the passive layer was no longer

be able to form on the sample. The Eb is shown to be similar for all four above mentioned alloys.

These results further demonstrate the improved corrosion performance of this new low-Cu

content alloy as previously shown by Li et al. when comparing to Zr60Ni25Al15 and

Zr56Ni25Al15Nb4 (shown in Figure 6) [63]. Lastly, these results define the Eb which can be used

as a baseline for determine appropriate ECMM machining voltages.

63

Figure 30: Anodic polarization scan for Sample III in 2.94 M NaNO3 at room temperature and

open air with scan rate of 1 mV/s. Section I is the passive region, II is the O2 reaction, III is a

change is oxidation states and IV is the pitting region.

To further test the corrosion behaviour of this alloy system, a comparison between partially-

crystalline samples and amorphous ones was conducted. Figure 31 below is a graph comparing

the anodic polarization scans between an amorphous sample (V) and crystalline (IV) one. Both

scans were produced using a scan speed of 1 mV/s in 2.94 M NaNO3 at room temperature and

open to air. Appendix D shows duplicate scans for each sample and validates the data. From

Figure 31, it can be seen that the partially-crystalline sample has a slightly higher Ecorr than the

amorphous sample and is in contrast to what would be expected. XRD data for Sample IV

showed the presence of some crystalline oxide phases which could help to explain the higher

Ecorr compared to the amorphous sample. In order to confirm these Ecorr values are accurate and

that partially-crystalline samples do show a higher value than amorphous samples, Tafel

measurements would need to be conducted. Tafel measurements would allow for more accurate

measurements of Ecorr. Although the Ecorr values are very similar between the two samples, it is

1.00E-10

1.00E-09

1.00E-08

1.00E-07

1.00E-06

1.00E-05

1.00E-04

1.00E-03

1.00E-02

1.00E-01

1.00E+00

-1 -0.5 0 0.5 1 1.5 2 2.5

Cu

rren

t D

ensi

ty, I

(A

/cm

2)

Potential, E (VAg/AgCl)

2.94 M NaNO3 RT, open air

Eb

I II III IV

64

clear that Eb is higher for the amorphous sample (2.17 ± 0.01 V for amorphous vs. 2.10 ± 0.03 V

for crystalline). Additionally, the Eb for amorphous samples is seen to be more reproducible

compared to crystalline samples. This observation is due to the amorphous samples having a

chemically homogenous structure as where the distribution of crystalline phases can be random

in the crystalline samples. This result is consistent with previous reports and suggests the

amorphous structure is more resistant against pitting [65].

Figure 31: Anodic polarization comparison between amorphous and crystalline samples.

4.3.2 Chronoamperometry

While anodic polarization provides useful information about the corrosion behaviour of this alloy

and provides a general machining region, it should be noted that the scans are produced by

slowly increasing voltage in one direction. The machining process involves pulsing a high

enough voltage to cause a pit and then removing the applied voltage. As such, when the applied

voltage is removed, the sample will repassivate and be in a different part of the anodic

polarization scans. Therefore, to obtain a more accurate measurement of the machining region

1.E-09

1.E-08

1.E-07

1.E-06

1.E-05

1.E-04

1.E-03

1.E-02

1.E-01

1.E+00

1.E+01

-1 -0.5 0 0.5 1 1.5 2 2.5

Cu

rren

t D

ensi

ty, I

(A

/cm

2)

Potential, E (V vs. SCE)

Amorphous

Crystalline

2.94 M NaNO3 RT, open air

65

the ‘electrochemical scratch’ technique detailed in Section 3.6.2 is applied. During this process it

was noted that at a base voltage of 1.785 V the current did not drop back down to 0 immediately

(seen in Figure 32). It wasn’t till a base voltage of 2.235 V was applied that the voltage never

returned to 0, indicating this potential to be the Erepass (see Figure 33). Comparing these results to

anodic polarization indicates these two base voltages both correspond fairly closely to the

breakdown of water and the Eb. While these values show similar results to those obtained during

anodic polarization, the Erepass is a better value for use in ECMM.

Figure 32: Current response to 1.785 V base during chronoamperometry.

Figure 33: Current response to 2.235 V base during chronoamperometry.

0 50 100 150 200 250 300 350

Cu

rre

nt

(arb

. un

its)

Time (s)

Base - 1.785 V

0 50 100 150 200 250 300 350

Cu

rre

nt

(arb

. un

its)

Time (s)

Base - 2.235 V

66

4.4 Electrochemical Micromachining

The first attempt to perform ECMM on Zr-based BMGs was done in 2011 by Koza et al. [73].

He concluded that using an aqueous based electrolyte, such as NaNO3, for ECMM was not

possible due to the formation of thick and dense corrosion products. Figure 34 below are ECMM

holes made using 3 V and pulsing at 10 µs on: 100 µs off. It should be noted that 10 µs is the

fastest the potentiostat can pulse. Additionally, since the working distance could not be directly

controlled or measured, the aspect ratio is a better indicator of the size of the features. While no

signs of corrosion products were viewed on the surface after rinsing with acetone and distilled

water, the features observed were irregular in size and pitted. In contrast, Figure 35 shows a

machined hole when pulsing is not applied. It can be seen that the machining becomes much

more irregular and non-uniform. As such, through literature, initial attempts and

chronoamperometry, it was determined that the peak to base voltage, electrolyte flow rate and

pulsing were important parameters on the ECMM process. Table 9 in Section 3.6.3 outlines the

various parameters that were tested, while Table 10 outlines the experiments done to assess

them.

Figure 34: SEM image of ECMM hole at 3 V 10µs on: 100 µs off, no electrolyte jet.

67

Figure 35: SEM image of ECMM hole at 3V, no pulse, no electrolyte jet.

After performing each of the experiments in Table 9, the aspect ratio, roughness within holes,

and depth of holes was tabulated and is presented in Table 17. The corresponding SEM images

can be viewed in Figure 36 below along with the corresponding parameters used. Duplicates of

each sample were prepared to generate an error range for the aspect ratio. Errors associated with

surface roughness and depths are given by the root mean squared (RMS) value. It should be

noted that due to some holes being very small in size, the surface roughness measurement picked

up some signal from the surface as well as the holes itself. For Run 3, the depth of the hole was

beyond the measuring range of profilometry (10 µm) and as a result the exact depth has not been

recorded. To measure this depth, micron probes can be used or taking a mould replica which will

be investigated in the future. These results are further discussed below.

68

Table 18: List of aspect ratios, surface roughness, depth, RMS and comments for ECMM holes.

Run Aspect Ratio Ra (μm) Depth (μm) RMS Comment

3V

Sample

1 1.41 ± 0.17 0.77 2.478 0.9 Irregular

2 1.25 ± 0.18 1.658 6.281 1.65 Stepped-

circular

3 1.29 ± 0.07 0.091 >10 0.109 Circular

4V

Sample

4 1.44 ± 0.31 0.378 1.278 0.422 Irregular

5 1.59 ± 0.58 0.592 2.034 0.635 Oval

6 1.63 ± 0.37 0.308 1.264 0.35 Oval

5V

Sample

7 1.04 ± N/A 0.134 1.074 0.218 Irregular

8 1.32 ± 0.21 0.118 0.811 0.166 Stepped-oval

9 1.37 ± 0.30 0.099 1.167 0.179 Oval

69

Figure 36: SEM micrographs for ECMM holes. Numbers correspond to experiment number in Table 9.

70

As previously stated, the working distance was not able to be fixed and determined from sample

to sample. Therefore to study the effects of the various parameters, each sample was analyzed

individually and then compared overall. The results showed that having a high duty cycle (i.e.

1:10) produced the best roughness, a high flow rate gave the greatest depths. The best aspect

ratio was achieved when using 2.235 V as the base voltage. A high duty cycle allows for

efficient use of electrons in order to machine away high energy areas such as peaks that are left

behind from the previous pulses. Pulsing too fast may not give enough time for these to form and

alter the feature size. The use of flow rate was to remove any corrosion products that may adhere

on the surface and therefore removal should also result in the greatest depths as well. Lastly, the

use of a base voltage of 2.235 V was used such that the area to be machined would not

spontaneously repassivate. Without having to break through a passive layer on each pulse, the

holes are able to be more uniform in diameter. This result is in contrast with previous results seen

when a constant voltage is applied and thus more analysis into the combination of this base

voltage with other parameters needs to be investigated. Even though 5V samples showed low

surface roughness overall, it was seen in SEM that pitting could be observed within these

features. These results are consistent with observations for ECMM of Fe-based BMGs by Horn

et al. [74].

Overall, the best combination of aspect ratio, roughness and depths was found in experiment

number 3. This experiment used a voltage range of 2.235 – 3 V with 1:10 duty cycle and 0.4

l/min flow rate. These parameters were then used to test the feasibility of producing lines on a

surface in attempt to be able to pattern. Figure 37 shows line produced using these parameters at

a movement speed of 0.22 µm/s. It can be seen that even though a high flow rate was used,

corrosion products can be found on the top line. The depths of these lines are not very deep and

can attribute to the somewhat fast movement speed. Horn et al. [74] was able to produce

complex shapes in Fe-based BMGs using a movement speed of 0.1 µm/s, which is significantly

slower. Pairing this with their use nano pulses, the features seen here are not as refined. With that

said, Figure 37 does show there is a good dimensional tolerance and accuracy associated with

this technique. With further improvements in electrolyte, faster pulsing and slower movement

speeds, the same type of features should be possible in Zr-based BMGs.

71

Figure 37: ECMM lines produce using experiment 3 parameters at 0.22 µm/s.

4.5 Abrasive Water Jet Machining

Although abrasive water jet machining (AWJM) has been shown to be a promising technique for

machining BMGs, the dimensional tolerance and surface roughness leave much to be desired

[75]. Abrasive water jet polishing can be performed in order to improve these two features;

however, this takes away from the benefit – time. Being able to determine processing conditions

that would allow precise, accurate and fast cutting are required to make AWJM a viable

technique for machining BMGs. Even though speed is not a key factor for machining the flexure

material, it is beneficial for producing and shaping test specimens for mechanical testing.

Additionally, there is no drawback if AWJM can be used without polishing to produce an

appropriate surface finish and tolerance in less time. Figure 38 below is the top surface of an

amorphous BMG cut using AWJM. Near the edges the sample is particularly rough and shows

signs of material flow during machining. These images correspond to Trial 1 in Table 11.

72

Figure 38: Trial 1 top surface of AWJM cut with magnification of edge on right.

Even though the dimensional tolerance and roughness are poor for this sample, XRD was

performed to verify no crystallization occurred during the machining process. Figure 39 below

shows the XRD pattern that was obtained after AWJM was performed. The pattern shows a

broad peak with no clear crystalline peaks present. To further confirm these results, the cross-

section of the sample was investigated under SEM. No signs of crystallization could be observed

in SEM, indicating that AWJM did not induce crystallization.

Figure 39: XRD pattern for Zr-Ni-Al-Cu-Nb BMG after Trial 1 AWJM process.

73

Figure 40: SEM cross section images after Trial 1 AWJM. Bulk (left) shows poor dimensional

tolerance as noted by cut line. Right image is a magnified image of the circle region.

Since AWJM has been confirmed to not cause crystallization, the focus shifted towards

improving the dimensional tolerance and surface roughness. Trial 2, from Table 11, focused on

the use of finer abrasive and placing the rod in sacrificial aluminum to eliminate the defects at

the edges seen in Figure 38 and 40. The effects of these parameter changes can be viewed in

Figures 41 and 42. Figure 41 shows the bulk top surface and magnification of centre features and

Figure 42 shows the magnified surface features and the entry and exit point of the water jet. The

cross section can be viewed in Figure 43 and was used to determine the dimensional tolerance.

Figure 41: Trial 2 AWJM top surface and along with magnification of surface features in centre.

74

Figure 42: Trial 2 exit (left) and entry (right) SEM images of top surface. Arrows indicate

direction of water jet machining.

Figure 43: Cross section image of Trial 2 cut showing very good dimensional tolerance across

surface.

75

Results from Trial 2 show huge improvements in the dimensional tolerance when using more

fine abrasive and sacrificial material. In fact, the dimensional tolerance was found to be ± 0.016

mm over the diameter of the rod. Surface roughness measurements revealed an average

roughness of 1.68 µm. While the dimensional tolerance is slightly above zero, there were

instances where there was no observed change in dimensions across the surface. Additionally,

the surface roughness achieved through AWJM meets the original criteria for < 2µm. Therefore,

it can be concluded that AWJM is a very useful technique in the machining and shaping of

BMGs with high speeds and excellent dimensional tolerance and surface roughness provided the

correct operating conditions are used.

76

5.0 Conclusions

In this work, it was shown that (Zr56Ni20Al15Cu5)100-xNbx bulk metallic glass alloys could be

successfully and repeatedly produced. This alloy contains the required mechanical and chemical

properties required as a flexure material in AGG. The crystallization behaviour, glass forming

ability and electrochemical behaviour were investigated. Machining using ECMM and AWJM

proved to be viable techniques. The following conclusions from this investigation can be drawn:

i) A high glass forming range was discovered between 2 and 4 at.% Nb. Alloys containing

Nb concentrations slightly above or below this range were found to have formed

Zr3NiO1.17 and CuZr2 upon casting.

ii) Increased oxygen content in the form of solute element destabilized the supercooled

liquid region and formed crystals upon casting, even within the high glass forming range.

Pre-existing oxide crystals also caused crystallization, but were much less detrimental on

the GFA. Therefore, it was determined that not only does the amount of oxygen present

play a key role on the GFA, but also the form in which it exists within BMG alloys.

ii) Crystallization behaviour was studied by deconvoluting multiple peaks seen during DSC

heating of amorphous samples. Using XRD, the each peak was identified. It was

determined that the first crystallization peak corresponded to c-Zr2Ni formation which

then transformed into either t-Zr2Ni or Zr5Ni4Al.

iv) Thermal analysis for amorphous samples using supercooled liquid region and activation

energy showed the 4at% Nb alloy to have higher values than the 2.7% Nb alloy. This

result suggests the critical casting diameter for this alloy system could be even further

increased using higher Nb content.

v) Anodic polarization and chronoamperometry were used to study the corrosion and pitting

behaviour of this Zr-based BMG in 2.94 M NaNO3. Chronoamperometry determined

Erepass to be 2.235 V. Using this value as the base voltage during ECMM led to more

uniform machining features.

77

vi) ECMM proved to be a viable micron-machining technique for Zr-based BMGs without

causing crystallization to occur. The use of an electrolyte jet to remove corrosion product

from the surface improved the depths of machining. Surface roughness was also reduced

when using a 1:10 on to off time during machining.

vii) AWJM was also proven to be an effective method for machining BMGs without

crystallization. Results showed that when a sacrificial piece of metal and fine abrasive are

used, excellent surface finishes of ± 0.016 mm and roughness of 0.17 µm can be

obtained.

78

6.0 Future Work

Fabrication of a BMG with desired mechanical properties has been conducted in this work;

however, more mechanical testing needs to be conducted to ascertain their utility in room

temperature AGG application. Zr-based BMGs are known to have very low loss coefficients, but

proper testing and identification is required. Development of an apparatus capable of sensing loss

coefficients is required. In addition, proper shaping of BMG into test component is required

since direct casting is not possible.

Both ECMM and AWJM were shown to be viable machining techniques at different depth scales

without causing crystallization to occur and are capable of achieving desired surface finish.

Various machining parameters have been analyzed in ECMM, but could be further improved

with selecting a more appropriate electrolyte. Additionally, better control over the working

distance will allow for improved and reproducible machining. While AWJM has shown very

promising results, a more systematic approach could be implemented to test the various

parameters used during machining.

79

References

[1] S. J. Thorpe, P. Borges, S. Shankar, Y. Li, X. Zeng, D. Cheung, K. A. Carroll and B.

French, "Application of Glassy Metals to Airborne Gravity Gradiometry," Ontario Centres

of Excellence.

[2] D. Turnbull, "Metastable structures in metallurgy," Metallurgical Transactions B, vol. 12,

no. 5, pp. 695-708, 1981.

[3] D. Turnbull, "Under what conditions can a glass be formed?," Contemporary physics, vol.

10, pp. 473-488, 1969.

[4] Q. K. Jiang, X. D. Wang, X. P. Nie, G. Q. Zhang, H. Ma, H. -J. Fecht, J. Bendnarcik, H.

Franz, Y. G. Liu, Q. P. Cao and J. Z. Jiang, "Zr-(Cu,Ag)-Al bulk metallic glasses," Acta

Materialia, vol. 56, pp. 1785-1796, 2008.

[5] H. A. Davies, "The kinetics of formation of a Au-Ge-Si metallic glass," Journal of Non-

Crystalline Solids, vol. 17, no. 2, pp. 266-272, 1975.

[6] H. Men, Z. Q. Hu and J. Xu, "Bulk metallic glass formation in the Mg-Cu-Zn-Y system,"

Scripta Materialia, vol. 46, pp. 699-703, 2002.

[7] X. Hu, S. C. Ng, Y. P. Feng and Y. Li, "Glass forming ability and in-situ composite

formation in Pd-based bulk metallic glass," Acta Materialia, vol. 51, pp. 561-572, 2003.

[8] W. H. Wang, C. Dong and C. H. Shek, "Bulk metallic glasses," Materials Science and

Engineering R, vol. 44, pp. 45-89, 2004.

[9] A. Inoue, T. Zhang and N. Nishiyama, "Preparation of 16 mm diameter rod of amorphous

Zr65Al7.5Ni10Cu17.5," Materials Transaction, vol. 34, no. 12, pp. 1234-1237, 1993.

[10] A. Peker and W. L. Johnson, "A highly processable metallic glass:

Zr41.2Ti13.8Cu12.5Ni10Be22.5".

[11] C. Suryanarayana and A. Inoue, "Metallic Glasses," in Bulk Metallic Glasses, Boca Raton,

CRC Press, 2011, p. 38.

[12] A. L. Greer, "Confusion by design," Nature, vol. 366, pp. 303-304, 1993.

[13] A. Inoue, "Stabilization of metallic supercooled liquid and bulk amorphous alloys," Acta

Materialia, vol. 48, no. 1, pp. 279-306, 2000.

80

[14] W. H. Wang, "Roles of minor additions in formation and properties of bulk metallic

glasses," Progress in Materials Science, vol. 52, pp. 540-596, 2007.

[15] H. K. Bhadeshia and R. W. Honeycombe, Steels: microstructure and properties, 3rd edn.,

London, U.K.: Butterworths-Heinemann, 2006.

[16] W. Zhang, Q. Zhang, C. Qin and A. Inoue, "Synthesis and properties of Cu-Zr-Ag-Al glassy

alloys with high glass-forming ability," Materials Science & Engineering B, vol. 148, pp.

92-96, 2008.

[17] J. F. Löffler, "Bulk metallic glasses," Intermetallics, vol. 11, pp. 529-540, 2003.

[18] W. Zhang, Q. Zhang and A. Inoue, "Formation and thermal stability of new Zr-Cu-based

bulk glassy alloys with unusual glass-forming ability," Journal of Alloys and Compounds,

vol. 483, pp. 112-115, 2009.

[19] X. Wang, Q. P. Cao, Y. M. Chen, K. Hono, C. Zhong, Q. K. Jiang, X. P. Nie, L. Y. Chen,

X. D. Wang and J. Z. Jiang, "A plastic Zr-Cu-Ag-Al bulk metallic glass," Acta Materialia,

vol. 2011, no. 59, pp. 1037-1047, 2011.

[20] H. E. Kissinger, "Reaction kinetics in differential thermal analysis," Analytical Chemistry,

vol. 29, no. 11, pp. 1702-1706, 1957.

[21] T. Ozawa, "Kinetic analysis of derivative curves in thermal analysis," Journal of Thermal

Analysis, vol. 2, pp. 301-324, 1970.

[22] M. Avrami, "Kinetics of Phase Change. I General Theory," Journal of Chemical Physics,

vol. 7, pp. 1103-1112, 1939.

[23] M. Avrami, "Kinetics of Phase Change. II Transformation‐Time Relations for Random

Distribution of Nuclei," Journal of Chemical Physics, vol. 8, pp. 212-224, 1940.

[24] C. R. Zhou, K. Lu and J. Xu, "The effect of nitride addition on the thermal stability of

mechanically alloyed amorphous Zr65Al7.5Cu17.5Ni10 alloy," Journal of Metastable and

Nanocrystalline Materials, vol. 8, pp. 116-122, 2000.

[25] W. L. Johnson, "Bulk Glass-Forming Metallic Alloys: Science and Technology," MRS

Bulletin, vol. 24, no. 10, pp. 42-56, 1999.

[26] A. Inoue, D. Kawase, A. P. Tsai, T. Zhang and T. Masumoto, "Stability and transformation

to crystalline phases of amorphous Zr65Al7.5Cu27.5 alloys with significant supercooled

81

liquid region," Materials Science and Engineering A, vol. 178, no. 1-2, pp. 255-263, 1994.

[27] J. Schroers, B. Lohwongwatana, W. L. Johnson and A. Peker, "Gold based bulk metallic

glass," Applied Physics Letters, vol. 87, pp. 061912-1-061912-3, 2005.

[28] C. Chou and D. Turnbull, "Transformation behavior of Pd-Au-Si metallic glasses," Journal

of Non-Crystalline Solids, vol. 17, pp. 169-188, 1975.

[29] F. Q. Guo, H. J. Wang, S. J. Poon and G. J. Shiflet, "Ductile titanium-based glassy alloy

ingots," Applied Physics Letters, vol. 86, pp. 091907-1-091907-3, 2005.

[30] F. Q. Guo, S. J. Poon and G. J. Shiflet, "CaAl-based bulk metallic glasses with high thermal

stability," Applied Physcis Letters, vol. 84, pp. 37-39, 2004.

[31] A. Inoue, N. Nishiyama and H. M. Kimura, "Preparation and thermal stability of bulk

amorphous Pd40Cu30Ni10P20 alloy cylinder of 72mm in diameter," Material Transactions,

vol. 38, pp. 179-183, 1997.

[32] A. Inoue, T. Zhang, N. Nishiyama, K. Ohba and T. Masumoto, "Preparation of 16mm

diameter rod of amorphous Zr65Al7.5Ni10Cu17.5 alloy," Material Transactions, vol. 34,

pp. 1234-1237, 1993.

[33] A. Inoue, T. Shibata and T. Zhang, "Effect of Additional Elements on Glass Transition

Behavior and Glass Formation Tendency of Zr-Al-Cu-Ni Alloys," Materials Transactions,

vol. 36, no. 12, pp. 1420-1426, 1995.

[34] W. Dong, H. Zhang, W. Sun, B. Ding and Z. Hu, "Formation, Thermal Stability and

Mechanical Properties of Zr-Nb-Cu-Ni-Al Bulk Metallic Glasses," Materials Transactions,

vol. 47, no. 5, pp. 1294-1298, 2006.

[35] A. Inoue, T. Zhang and T. Masumoto, "Preparation of bulky amorphous Zr-Al-Co-Ni-Cu

alloys by copper mold casting and their thermal and mechanical properties," Materials

Transactions, vol. 36, no. 3, pp. 391-398, 1995.

[36] T. Zhang and A. Inoue, "Formation, thermal and mechanical properties of bulk glassy alloys

in Zr–Al–Co and Zr–Al–Co–Cu systems," Materials Science and Engineering: A, Vols.

375-377, pp. 432-435, 2004.

[37] X. Xiao, S. Fang, L. Xia, W. Li, Q. Hua and Y. Dong, "Thermal and mechanical properties

of Zr52.5Al10Ni10Cu15Be12.5 bulk metallic glass," Journal of Alloys and Compounds,

82

vol. 351, no. 1-2, pp. 324-328, 2003.

[38] A. Inoue, T. Zhang and T. Masumoto, "Zr-Al-Ni amorphous alloys with high glass

transition temperature and significant supercooled liquid region," Materials Transactions,

vol. 31, no. 3, pp. 177-183, 1990.

[39] H. Kakiuchi, A. Inoue, M. Onuki, Y. Takano and T. Yamaguchi, "Application of Zr-based

bulk glassy alloys to golf clubs," Materials Transactions, vol. 42, no. 4, pp. 678-681, 2001.

[40] Y. Yokoyama, K. Fukaura and A. Inoue, "Formation and mechanical properties of Zr-Cu-Al

bulk glassy alloys," Materials Science and Engineering: A, Vols. 375-377, pp. 427-431,

2004.

[41] H. Kato, J. Saida and A. Inoue, "Influence of hydrostatic pressure during casting on as cast

structure and mechanical properties in Zr65Al7.5Ni10Cu17.5−xPdx (x = 0, 17.5) alloys,"

Scripta Materialia, vol. 51, no. 11, pp. 1063-1068, 2004.

[42] H. Kimura, M. Kishida and T. Kaneko, "Physical and mechanical properties of Zr-based

metallic glass," Materials Transactions, vol. 36, no. 7, pp. 890-895, 1995.

[43] T. Zhang and A. Inoue, "Density, thermal stability and mechanical properties of Zr-Ti-Al-

Cu-Ni bulk amorphous alloys with high Al plus Ti concentrations," Material Transaction,

vol. 39, no. 8, pp. 857-862, 1998.

[44] T. Zhang and A. Inoue, "Mechanical properties of Zr-Ti-Al-Ni-Cu bulk amorphous sheets

prepared by squeeze casting," Materials Transaction, vol. 39, no. 12, pp. 1230-1237, 1998.

[45] K. Jin and J. F. Loffler, "Bulk metallic glass formation in Zr–Cu–Fe–Al alloys," Applied

Physics Letters, vol. 86, p. 241909, 2005.

[46] X. D. Hui, H. C. Kou, J. P. He, Y. L. Wang, W. Dong and G. L. Chen, "Preparation,

microstructure and mechanical properties of Zr-based bulk amorphous alloys containing

tungsten," Intermetallics, vol. 10, no. 11-12, pp. 1065-1069, 2002.

[47] K. Q. Qiu, A. M. Wang, H. F. Zhang, B. Z. Ding and Z. Q. Hu, "Mechanical properties of

tungsten fiber reinforced ZrAlNiCuSi metallic glass matrix composite," Intermetallics, vol.

10, no. 11-12, pp. 1283-1288, 2002.

[48] R. D. Conner, A. J. Rosakis, W. L. Johnson and D. M. Owen, "Fracture toughness

determination for a beryllium-bearing bulk metallic glass," Scripta Materialia, vol. 37, no.

83

9, pp. 1373-1378, 1997.

[49] Y. Zhang, D. Q. Zhao, B. C. Wei, P. Wen, M. X. Pan and W. H. Wang, "Formation and

properties of Zr48Nb8Fe8Cu12Be24 bulk metallic glass," Journal of Materials Research,

vol. 16, no. 06, pp. 1675-1679, 2001.

[50] A. Inoue, C. Fan, J. Saida and T. Zhang, "High-strength Zr-based bulk amorphous alloys

containing nanocrystalline and nanoquasicrystalline particles," Science and Technology of

Advanced Materials, vol. 1, no. 2, pp. 73-86, 2000.

[51] L. Q. Xing, G. P. Gorier and D. M. Herlach, "Cast bulk Zr57Ti5Al10Cu20Ni8 amorphous

alloy with tendency of phase separation," Materials Science and Engineering: A, Vols. 226-

228, pp. 429-433, 1997.

[52] C. T. Liu, M. F. Chisholm and M. K. Miller, "Oxygen impurity and microalloying effect in

a Zr-based bulk metallic glass alloy," Intermetallics, vol. 10, no. 11-12, pp. 1105-1112,

2002.

[53] Z. Bian, M. X. Pan, Y. Zhang and W. Hua, "Carbon-nanotube-reinforced

Zr52.5Cu17.9Ni14.6Al10Ti5Zr52.5Cu17.9Ni14.6Al10Ti5 bulk metallic glass composites,"

Applied Physics Letters, vol. 81, p. 4739, 2002.

[54] G. He, W. Loser, J. Eckert and L. Schultz, "Phase transformation and mechanical properties

of Zr-base bulk glass-forming alloys," Materials Science and Engineering: A, vol. 352, no.

1-2, pp. 179-185, 2003.

[55] J. Eckert, U. Kuhn, J. Das, S. Scudino and N. Radtke, "Nanosctructured composite materials

with improved deformation behavior," Advanced Engineering Materials, vol. 7, no. 7, pp.

587-596, 2005.

[56] Y. F. Sun, S. K. Guan, B. C. Wei, Y. R. Wang and C. H. Shek, "Brittleness of Zr-based bulk

metallic glass matrix composites containing ductile dendritic phase," Materials Science and

Engineering: A, vol. 406, no. 1-2, pp. 57-62, 2005.

[57] X. F. Pan, H. Zhang, Z. F. Zhang, M. Stoica, G. He and J. Eckert, "Vickers hardness and

compressive properties of bulk metallic glasses and nanostructure-dendrite composites,"

Journal of Materials Research, vol. 20, no. 10, pp. 2632-2638, 2005.

[58] A. Inoue, T. Zhang, M. W. Chen and T. Sakurai, "Mechanical Properties of Bulk

84

Amorphous Zr-Al-Cu-Ni-Ag Alloys Containing Nanoscale Quasicrystalline Particles,"

Material Transaction, vol. 40, no. 12, pp. 1382-1389, 1999.

[59] T. C. Hufnagel, T. Jiao, Y. Li, L.-Q. Xing and K. T. Ramesh, "Deformation and failure of

Zr57Ti5Cu20Ni8Al10 bulk metallic glass under quasi-static and dynamic compression,"

Journal of Materials Research, vol. 17, no. 06, pp. 1441-1445, 2002.

[60] X. Rao, P. C. Si, J. N. Wang, Z. Xu, S. Xu, W. M. Wang and W. H. Wang, "Preparation and

mechanical properties of a new Zr–Al–Ti–Cu–Ni–Be bulk metallic glass," Materials

Letters, vol. 50, no. 5-6, pp. 279-283, 2001.

[61] X. Gu, T. Jiao, L. J. Kecskes, R. H. Woodman, C. Fan, K. T. Ramesh and T. C. Hufnagel,

"Crystallization and mechanical behavior of (Hf, Zr)–Ti–Cu–Ni–Al metallic glasses,"

Journal of Non-Crystalline Solids, vol. 317, no. 1-2, pp. 112-117, 2003.

[62] A. Inoue and W. Zhang, "Formation, thermal stability and mechanial properties of Cu-Zr

and Cu-Hf binary glassy alloy rods," Materials Transactions, vol. 45, no. 2, pp. 584-587,

2004.

[63] Y. Li, W. Zhang, F. Qin and A. Makino, "Mechanical properties and corrosion resistance of

a new Zr56Ni20Al15Nb4Cu5 bulk metallic glass with a diameter up to 25 mm," Journal of

Alloys and Compounds, vol. 615, pp. 571-574, 2014.

[64] W. H. Peter, C. T. Buchanan, P. K. Liu, M. L. Morrison, J. A. Horton, C. A. Carmichael Jr.

and J. L. Wright, "Localized corrosion behaviour of a zirconium-based bulk metallic glass

relative to its crystalline state," Intermetallics, vol. 10, pp. 1157-1162, 2002.

[65] V. Schroeder, C. j. Gilbert and R. O. Ritchie, "Comparison of the corrosionbehavior of a

bulk amorphous metal, Zr41.2Ti13.8Cu12.5Ni10Be22.5, with its crystallized form," Scripta

Materialia, vol. 38, pp. 1481-1485, 1998.

[66] L. Liu, C. L. Qiu, H. Zou and K. C. Chan, "The effect of the microalloying of Hf on the

corrosion behavior of ZrCuNiAl bulk metallic glass," Journal of Alloys and Compounds,

vol. 399, no. 1-2, pp. 144-148, 2005.

[67] K. Asami, H. Habazaki, A. Inoue and K. Hashimoto, "Recent development of highly

corrosion resistant bulk glassy alloys," Materials Science Forum, vol. 502, pp. 225-230,

2005.

85

[68] J. Igel, "Production of Bulk Metallic Glasses for Use in Airborne Gravity Gradiometry,"

University of Toronto, Toronto, 2016.

[69] H. A. Youssef and H. El-Hofy, Machining technology : machine tools and operations, Boca

Raton, Fl: Taylor & Francis, 2008.

[70] A. Spieser and A. Ivanov, "Recent developments and research challenges in electrochemical

micromachining (uECM)," International Journal of Advanced Manufacturing Technology,

vol. 69, pp. 563-581, 2013.

[71] P. Loc and F.-J. Shiou, "Abrasive water jet polishing on Zr-based bulk metallic glass,"

Advanced Materials Research, vol. 579, pp. 211-218, 2012.

[72] B. Bhattacharyya, J. Munda and M. Malapati, "Advancement in electrochemical micro-

machining," Machine tools & Manufacture, vol. 44, pp. 1577-1589, 2004.

[73] J. A. Koza, R. Sueptitz, M. Uhlemann, L. Schultz and A. Gebert, "Electrochemical

micromachining of a Zr-based bulk metallic glass using micro-tool electrode technique,"

Intermetallics, vol. 19, pp. 437-444, 2011.

[74] S. Horn, S. Oswald, M. Stoica, M. Uhlemann and A. Gebert, "Complex microshaping of

bulk metallic glass surfaces by electrochemical means," Journal of Materials Research, vol.

22, pp. 3493-3501, 2015.

[75] V. Wessels, A. Grigoryev, C. Dold, C.-F. Wyen, R. Roth, E. Weingartner, F. Pude, K.

Wegener and J. F. Loffler, "Abrasive waterjet machining of three-dimensional structures

from bulk metallic glasses and comparison with other techniques," Journal of Materials

Research, vol. 27, pp. 1187-1192, 2012.

[76] ASTM International, ASTM E1409: Standard Test Method for Determination of Oxygen and

Nitrogen in Titanium and Titanium Alloys by Inert Gas Fusion, 2013.

[77] A. Gebert, P. F. Gostin, R. Sueptitz, S. Oswald, S. Abdi, M. Uhlemann and J. Eckert,

"Polarization studies of Zr-based bulk metallic glasses for electrochemical machining,"

Journal of The Electrochemical Society, vol. 161, no. 4, pp. E66-E73, 2014.

[78] R. Baboian, Corrosion Tests and Standards: Application and Interpretation - Second

Edition, Baltimore, MD: ASTM International, 2005.

[79] CH Instruments, "Accessories," CH Instruments, 2016. [Online]. Available:

86

http://www.chinstruments.com/accessories.shtml#chi116_image. [Accessed 3 May 2016].

[80] Z. Zhu, W. Zhang, G. Xie and A. Inoue, "Relation between glass and quasicrystal formation

in the Zr–Nb–Cu–Ni–Al alloys upon solidification," Applied Physics Letters, vol. 97, p.

031919, 2010.

[81] C. Suryanarayana and A. Inoue, "Computational Approaches," in Bulk Metallic Glasses,

Boca Raton, FL, CRC Press, 2011, pp. 126-128.

[82] A. Gebert, J. Eckert and L. Schultz, "Effect of oxygen on phase formation and thermal

stability of slowly cooled Zr65Al7.5Cu17.5Ni10 metallic glass," Acta Materialia, vol. 46,

no. 15, pp. 5475-5482, 1998.

[83] J. Eckert, N. Mattern, M. Zinkevitch and M. Seidel, "Crystallization behavior and phase

formation in Zr-Al-Cu-Ni metallic glass containing oxygen," Materials Transactions, JIM,

vol. 39, no. 6, pp. 623-632, 1998.

[84] B. S. Murty, D. H. Ping, K. Hono and A. Inoue, "Influence of oxygen on the crystallization

behavior of Zr65Cu27.5Al7.5 and Zr66.7Cu33.3 metallic glasses," Acta Materialia, vol. 48,

pp. 3985-3996, 2000.

[85] J. Igel, D. W. Kirk, C. V. Singh and S. J. Thorpe, "A Practical Investigation of the

Production of Zr-Cu-Al-Ni Bulk Metallic Glasses by Arc Melting and Suction Casting,"

Materials Transactions, vol. 56, no. 11, pp. 1834-1841, 2015.

[86] Y.-h. Li, W. Zhang, C. Dong, J.-b. Qiang and A. Makino, "Correlation between the glass-

forming ability and activation energy of crystallization for Zr75−xNi25Alx metallic

glasses," International Journal of Minerals, Metallurgy and Materials, vol. 20, no. 5, pp.

445-449, 2013.

[87] L. Lin, W. Zhi-Fang and C. Li, "A Kinetic Study of the Non-isothermal Crystallization of a

Zr-Based Bulk Metallic Glass," Chinese Physics Letters, vol. 19, no. 10, pp. 1483-1486,

2002.

[88] J. Pang, Y. Wang, M. J. Tan and K. M. Liew, "Thermal properties of metallic glasses:

Heating rate dependence and their correlation," Materials Letters, vol. 126, pp. 81-84, 2014.

[89] C. Suryanarayana and A. Inoue, "Thermal Stability of Metallic Glasses," in Bulk Metallic

Glasses, Boca Raton, FL, CRC Press Taylor & Francis Group, 2011, pp. 208-211.

87

[90] K. Cheng, "Evaluation of crystallization kinetics of glasses by non-isothermal analysis,"

Journal of Materials Science, vol. 36, pp. 1043-1048, 2001.

[91] H.-R. Wang, Y.-L. Gao, Y.-F. Ye, G.-H. Min, Y. Chen and X.-Y. Teng, "Crystallization

kinetics of an amorphous Zr-Cu-Ni alloy: calculation of the activation energy," Journal of

Alloys and Compounds, vol. 353, pp. 200-206, 2003.

[92] M. Iqbal and J. I. Akhter, "Influence of Y and Nb addition on Crystallization Behavior and

Mechanical Properties of Zr-Ni-Al-Cu-M Bulk Amorphous Alloys," Advanced Materials

Research, vol. 326, pp. 11-18, 2011.

[93] Y. X. Zhuang, T. F. Duan and H. Y. Shi, "Calorimetric study of non-isothermal

crystallization kinetics of Zr60Cu20Al10Ni10 bulk metallic glass," Journal of Alloys and

Compounds, vol. 509, pp. 9019-9025, 2011.

88

Appendix A: Detailed Casting Protocol for BMGs

The following Appendix is a detailed production process for casting BMGs by Josh Igel [68].

The following steps were performed in the production process:

1. All raw materials were weighed to the desired amount. Lineman’s pliers were used to cut

the slugs into smaller pieces when necessary to obtain the desired weight. Weighing was

performed with a Mettler AE260 DeltaRange® analytical balance with an accuracy of ±

0.2 mg. All raw materials were weighed to be within 0.4 mg of the desired weight.

2. The materials were transported to the arc melter in scintillation vials.

3. The arc melter main power switch (14) was switched on, the large nut (18) was

unscrewed and the base plate (7) was detached from the vacuum chamber (3).

4. The mould (19) was removed from the mould column (6) in the base plate (7) to grind the

inner cavity with 1200 grit SiC to obtain a smooth surface, and the base plate (7) was

cleaned with isopropanol. Further cleaning of the base plate (7) by grinding with 1200

grit SiC paper and polishing with 3 μm diamond paste was performed periodically when

deemed necessary to remove oxide build-up.

5. The mould (19) was returned to the mould column (6), and raw materials were loaded

into the crucible well (8) by ascending order of weight (i.e. the order was Al, Ni, Cu then

Zr) to avoid pieces from falling out of the crucible well (8) during this material loading

step. Loading material in ascending order of weight was done solely for ease of loading.

6. The base plate (7) was re-attached to the vacuum chamber (3). The bottom piece of the

mould (20) was placed in the middle of the large nut (18), and the large nut (18) was

89

screwed back on to the bottom of the base plate (7). If done correctly, the pin of the

bottom piece should fit into the cavity of the mould (19).

7. The main valve on the argon tank was opened and the regulator was set to deliver argon

at 250 kPa (based on the manufacturer’s specifications).

8. The vacuum pump power switch (13) was then switched on, and the vacuum pump

control knob (16) was turned counter-clockwise to the maximum position.

9. Once the pressure gauge (2) read -1.0 bar (which corresponds to approx. 0.1 mbar), the

vacuum pump control knob (16) was turned back to the zero position, and the vacuum

chamber (3) was filled with argon by opening the gas inlet valve (17). The gas inlet valve

(17) was closed when the pressure gauge (2) read 0 bar (corresponding to atmospheric

pressure).

10. Steps 8 and 9 were repeated three times. This pump down-argon filling process was

performed to avoid oxidation during melting by eliminating as much atmospheric oxygen

from the vacuum chamber (3) as practically possible.

11. The vacuum chamber (3) was evacuated again using the vacuum pump (15) as done in

Step 8. Once evacuated, the vacuum valve (9) was turned to the closed position to lock

low pressure argon (approx. 0.1 mbar) in the piping below the mould (19). The vacuum

pump power switch (13) was then switched off, and the vacuum chamber (3) was filled

with argon as done in Step 9. This was done to provide a pressure differential between

the top and bottom of the mould (19) to facilitate suction for casting.

12. The electrode power switch (11) was switched on, and the current level knob (12) was set

to 6.

90

13. A plasma arc was created between the base plate (7) and electrode (4) by touching the

electrode (4) to the igniter pin (5).

14. With a plasma arc created between the base plate (7) and electrode (4), the current level

knob (12) was reset to 3, and the arc was moved over the crucible well (8) containing the

raw materials. The arc was then circled over the crucible well to both melt and mix the

raw material. The current level knob (12), which controls the intensity of the arc, was set

to 3 to avoid material from being ejected from the crucible well (8) during melting. The

amount of time in which the plasma arc was circled over the raw material was dependent

on the particular homogenization protocol being used.

15. Once the melting time had elapsed, the electrode power switch (11) was switched off.

The alloy ingot was then flipped using the electrode (4) and re-melted. The number of

melting iterations was dependent on the homogenization protocol used.

16. Once homogenization had been completed based on the particular protocol being used,

the ingot was moved to the crucible at the top of the mould (19) and melted using the

electrode (4) as described earlier. Melting was performed until the ingot appeared

completely molten (approx. 4 seconds).

17. Once the ingot appeared completely molten, the vacuum valve (10) was turned to the

open position to release the pressure differential between the top and bottom of the mould

(19) causing the molten alloy to be drawn into the copper mould (19).

18. The vacuum chamber (3) was filled with argon to restore atmospheric pressure.

19. The large nut (18) was unscrewed, the base plate (7) was detached from the vacuum

chamber (3), and the solidified rod was extracted from the copper mould (19).

91

Figure 44: Detailed schematic of the Compact Arc Melter MAM-1 arc melter.

Figure 45: Image of vacuum chamber assembly and gas control panel.

92

Figure 46: Images of a) the copper base plate detached from the vacuum chamber and b) normal

view of the base plate.

Figure 47: Image of a) split copper mould and b) bottom piece.

93

Appendix B: XRD Database References

Table 19: XRD data for phases present during casting and heating of Zr-Ni-Al-Cu-Nb BMGs. The angle position (2θ), d-value (d),

intensity (i), and crystal plane (hkl) for the first three most intense peaks of each crystalline phase are given. ICSD card numbers are

given for reference.

Phase

Peak 1 Peak 2 Peak 3

Card number

2θ d i hkl 2θ d i hkl 2θ d i hkl

cubic-Zr2Ni 38.10 2.360 100 333 35.85 2.503 63.0 422 41.60 2.169 35.0 440 00-041-0898

tetragonal-Zr2Ni 35.29 2.541 100 211 44.27 2.044 31.7 202 44.11 2.052 29.5 310 03-065-0487

Zr5Ni4Al 39.00 2.307 100 212 39.75 2.266 42.1 310 36.85 2.437 31.7 202 01-075-5651

Zr2Cu 38.05 2.363 100 511 35.79 2.507 38.9 422 64.32 1.447 27.7 660 01-074-7476

Zr3NiO1.17 34.75 2.579 100 023 37.91 2.372 47.8 131 36.44 2.463 34.6 130 01-072-7452

m-ZrO2 28.18 3.164 100 -111 31.47 2.841 68.2 111 50.12 1.819 22.6 220 01-083-0939

t-ZrO2 30.34 2.944 100 101 34.68 2.584 28.4 002 60.46 1.530 22.9 211 01-070-7300

Nb0.96O0.04 38.29 2.349 100 110 69.23 1.356 24.1 211 55.27 1.661 14.1 200 01-074-6000

94

Appendix C: SEM Micrographs for As-cast BMG Rods

Figure 48: SEM micrographs for cross-sections for as-cast rods listed in Table 12 taken at 500x

magnification to compare crystalline features. Sample I, IV and VI show clear dendritic crystal

formation. Sample II, III and V do not show any indication of crystals; however, XRD indicates

oxides are present in Sample II that cannot be seen in SEM.

95

Appendix D: Kissinger Plots

Figure 49: Kissinger plot for Sample I (Table 13).

Figure 50: Kissinger plot for Sample II (Table 13).

y = -28.753x + 26.501 R² = 0.9993

-13

-12

-11

-10

-9

-8

1.24 1.26 1.28 1.3 1.32 1.34 1.36

ln (

T p2/β

)

1000/T (1/K)

Kissinger Plot - Sample I

y = -30.032x + 27.162 R² = 0.9935

-13

-12

-11

-10

-9

-8

1.22 1.24 1.26 1.28 1.3 1.32

ln (

T p2/β

)

1000/T (1/K)

Kissinger Plot - Sample II

96

Figure 51: Kissinger plot for Sample III (Table 13).

Figure 52: Kissinger plot for Sample IV (Table 13).

y = -39.036x + 39.501 R² = 0.9954

-13

-12

-11

-10

-9

-8

1.25 1.26 1.27 1.28 1.29 1.3 1.31 1.32 1.33

ln (

T p2/β

)

1000/T (1/K)

Kissinger Plot - Sample III

y = -36.801x + 35.218 R² = 0.9997

-13

-12

-11

-10

-9

-8

1.21 1.22 1.23 1.24 1.25 1.26 1.27 1.28 1.29 1.3

ln (

T p2/β

)

1000/T (1/K)

Kissinger Plot - Sample IV

97

Figure 53: Kissinger plot for Sample V (Table 13).

Figure 54: Kissinger plot for Sample VI (Table 13).

y = -36.943x + 36.608 R² = 0.9927

-13

-12

-11

-10

-9

-8

1.24 1.25 1.26 1.27 1.28 1.29 1.3 1.31 1.32 1.33

ln (

T p2/β

)

1000/T (1/K)

Kissinger Plot - Sample V

y = -32.495x + 30.079 R² = 0.997

-13

-12

-11

-10

-9

-8

1.2 1.22 1.24 1.26 1.28 1.3 1.32

ln (

T p2/β

)

1000/T (1/K)

Kissinger Plot - Sample VI

98

Appendix E: Ozawa Plots

Figure 55: Ozawa plot for Sample I (Table 13).

Figure 56: Ozawa plot for Sample II (Table 13).

y = -13.154x + 18.148 R² = 0.9994

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

1.24 1.26 1.28 1.3 1.32 1.34 1.36

log(

β)

1000/Tx (1/°C)

Ozawa Plot - Sample I

y = -13.725x + 18.456 R² = 0.9941

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

1.22 1.24 1.26 1.28 1.3 1.32

log(β

)

1000/Tx (1/K)

Ozawa Plot - Sample II

99

Figure 57: Ozawa plot for Sample III (Table 13).

Figure 58: Ozawa plot for Sample IV (Table 13).

y = -17.626x + 23.801 R² = 0.9957

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

1.25 1.26 1.27 1.28 1.29 1.3 1.31 1.32 1.33

log(β

)

1000/Tx (1/K)

Ozawa Plot - Sample III

y = -16.674x + 21.965 R² = 0.9997

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

1.21 1.22 1.23 1.24 1.25 1.26 1.27 1.28 1.29 1.3

log(β

)

1000/Tx (1/K)

Ozawa Plot - Sample IV

100

Figure 59: Ozawa plot for Sample V (Table 13).

Figure 60: Ozawa plot for Sample VI (Table 13).

y = -16.72x + 22.549 R² = 0.9933

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

1.24 1.25 1.26 1.27 1.28 1.29 1.3 1.31 1.32 1.33

log(β

)

1000/Tx (1/K)

Ozawa Plot - Sample V

y = -14.799x + 19.728 R² = 0.9972

0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

1.8

1.2 1.22 1.24 1.26 1.28 1.3 1.32

log(β

)

1000/Tx (1/K)

Ozawa Plot - Sample VI

101

Appendix F: Anodic Polarization

Figure 61: Anodic polarization scans for amorphous Sample V at a scan speed of 1 mV/s in 2.94

M NaNO3 at room temperature and open air.

1.00E-02

1.00E-01

1.00E+00

1.00E+01

1.00E+02

1.00E+03

1.00E+04

1.00E+05

1.00E+06

1.00E+07

1.00E+08

-1 0 1 2

Cu

rren

t D

ensi

ty, I

(A

/cm

2)

Potential, E (V vs. SCE)

Series2

Series1

Sample V - Amorphous 2.94 m NaNO3

Test 1 Test 2

102

Figure 62: Anodic polarization scans for partially-crystalline Sample IV at a scan speed of 1

mV/s in 2.94M NaNO3 at room temperature and open air.

1.00E-02

1.00E-01

1.00E+00

1.00E+01

1.00E+02

1.00E+03

1.00E+04

1.00E+05

1.00E+06

1.00E+07

1.00E+08

-1 -0.5 0 0.5 1 1.5 2 2.5

Cu

rren

t D

ensi

ty, I

(A

/cm

2)

Potential, E (V vs. SCE)

Series2

Series1

Sample IV - Crystalline 2.94 M NaNo3

Test 1 Test 2