CITY UNIVERSITY OF HONG KONG

DEPARTMENT OF

PHYSICS AND MATERIALS SCIENCE

BACHELOR OF ENGINEERING (HONS) IN MATERIALS ENGINEERING

2007-2008

DISSERTATION

Effects of a minor addition of Si, Sn and In on formation and mechanical properties of

Cu-Zr-Al bulk metallic glass

by

Lau Chung Yam

March 2008

Effects of a minor addition of Si, Sn and In on formation and mechanical properties of

Cu-Zr-Al bulk metallic glass

By

Lau Chung Yam

Submitted in partial fulfilment of the

requirements for the degree of

BACHELOR OF ENGINEERING (HONS)

IN

MATERIALS ENGINEERING

from

City University of Hong Kong

March 2008

Project Supervisor : Dr. C.H.Shek

Acknowledgements I would like to express my sincere gratitude to my project supervisor, Dr. C.H.Shek. His

patient guidance, invaluable suggestions and ideas throughout this project period were

deeply appreciated.

I would also like to thank for the help from my tutor, Mr. Jiliang Zhang for his

experienced technical support in so many areas. In addition, I would like to express my

appreciation to the staff, Mr T.F. Hung and Mr. K.C. Yuen who helped me a lot during

the SEM imaging and XRD examination.

I

Abstract

To develop bulk metallic glasses (BMGs) for structural applications, prevention of the

catastrophic failure caused by the formation and propagation of a single dominant shear

band is a must. As the bulk metallic glass formers showing stability with respect to

crystallization, many BMG matrix composites have been developed by making use of

various crystalline reinforcement phases to inhibit shear band propagation. However,

very few studies on the optimization of the microstructure and mechanical behavior of

these BMG matrix composites were done in the past.

In this thesis, Cu-Zr-Al metallic glass was selected to study the influence of 1 % atomic

ratio addition of Sn, Si, In, on the mechanical behaviors. The effect was studied at room

temperature (RT) using various techniques including differential scanning calorimetry

(DSC), X-ray diffraction (XRD), scanning electronic microscopy (SEM) etc. All glassy

alloys were prepared by arc melting and suction casting. Meanwhile, the effect of small

alloying addition on the thermal stability of Cu46Zr46Al8 was investigated.

XRD examination shows amorphous structures on all formulated glassy alloys. The

absent of sharp crystalline peak in diffraction pattern give evidence on the glassy forming

ability of Cu-Zr-Al-BMG with addition of chosen elements. In the finding of stiffness

dependence, Cu-Zr-Al-Sn BMG gave a remarkable increase in hardness in the Vickers

hardness test.; Cu-Zr-Al-Sn BMG with collapsed appearances were shown on the

indented edges. The occurrence of the collapsed edges can be explained by embrittlement

of amorphous matrix with nano-crystals. Both addition of Si, Sn to Cu-Zr-Al system

enhanced the stiffness of the referencing Cu-base BMG.

Study on the effect of silicon, tin, indium addition on the crystallization behavior of

Cu46Zr46Al8 was also performed. From the DSC curves obtained, an enlargement of the

supercooling region can be observed by alloying silicon together with the base Cu-Zr-Al

BMG, this implied that the thermal stability has been enhanced. The glass-forming ability

of Cu-Zr-Al metallic glass-forming alloys has been improved with 1% atomic addition of

II

III

silicon, which has been attributed to silicon destabilizing oxide nucleation sites and

increasing atomic size difference of the BMG.

Moreover, room temperature compression test revealed a significant strain hardening and

plastic strain elongation of 8.1% before failure for the 3-mm diameter (Cu46Zr46Al8)99Sn1

sample. The microstructure of the fracture surfaces of each parametric sample was

studied and compared under SEM. The remarkable enhancement on plastic deformation

can be explained in terms of shear band pattern and shear band density. They all give

evidences on the resistance of catastrophic failure of Cu-Zr-Al-Sn BMG.

List of Figures Figure2.1 Isothermal formation of a metastable metallic glass by interdiffusion at

temperature T d between two crystalline phases A and B

Figure 2.2

Simplified mechanism of arc melting suction casting method

Figure 2.3 Theoretical diagrams indicate different atomic arrangement of glassy

and crystalline structures upon shear force.

Figure 2.4 Typical strengths and elastic limits for various materials

Figure 2.5 Schematical atomic sizes of the chosen elements

Figure 3.2 DSC curve of the Cu46Zr46Al8 at heating rate of 20K/min

Figure 3.3 Three Portions of upper 2mm and lower 3mm diameter fabricated bulk

metallic glasses

Figure 4.1 DSC curves of the reference Cu46Zr46Al8 and with 1at % Sn, In, and Si

range from 400K to 580K

Figure 4.2 XRD patterns of Cu46Zr46Al8, (Cu46Zr46Al8)99Si1, (Cu46Zr46Al8)99Sn1

and (Cu46Zr46Al8)99In1 alloys.

Figure 4.3(a) Optical microscope photograph of 1kgf indentation of Cu-base BMG

with 1 at% Sn show the appearance of collapsed edges (Pointed by

arrow)

IV

Figure 4.3(b) Optical microscope photograph (with higher magnification) of 1kgf

indentation of Cu-base BMG with 1 at% Sn show the appearance of

collapsed edges (Pointed by arrow)

Figure 4.4(a) True stress–strain curves of upper part of (Cu46Zr46Al8)99Sn1 ingot

under uniaxial compression. Inset shows nominal stress–strain curves

of (Cu46Zr46Al8)99Sn1with three portions of the ingots

Figure 4.5(a) SEM images of the vein-like pattern of (Cu46Zr46Al8)99Sn1

Figure 4.5(b) SEM images of a site with higher dense of vein pattern at the core

Figure 4.6 SEM images show high dense shear bands with primary and secondary

shear bands. Shear offset shown in the inset image of sample Cu-Zr-

Al-Sn alloy

Figure 4.7(a) SEM of shear bands formed during uniaxial compression of

(Cu46Zr46Al8)99Si1. With arrows indicate the direction of secondary

SBs propagation

Figure 4.7(b) SEM of shear bands formed during uniaxial compression of

referencing Cu46Zr46Al8. With arrows indicate the direction of

secondary SBs propagation

V

VI

List of Tables

Table 2.1 Fundamental information of the chosen element

Table 3.1 Weight composition of different specimens

Table 4.1 Thermal parameters of (Cu46Zr46Al8)99X1 metallic glasses

Table 4.2 Summarize the calculated yield strength, young’s modulus, plastic strain

and hardness of all samples

Table of Content Acknowledgements…………………………………………………………… I

Abstract………………………………………………………………………… II

List of Figures ………………………………………………………………… IV

List of Tables…………………………………………………………………… VI

Table of Contents……………………………………………………………… VII

1. Introduction

2. Literature Review

2.1 Background of Metallic Glasses ………………………………… 2

2.2 Fabrication of Metallic Glasses ………………………………… 3

2.2.1 Critical cooling rate …………………………………… 3

2.2.2 Composite atomic size ………………………………… 3

2.2.3 Suction casting mechanism …………………………… 4

2.3 Mechanical Properties of Bulk Metallic Glasses ………………… 4

2.3.1 Compression Test at room temperature ………………… 4

2.3.2 Plastic deformation …………………………………… 4

2.4 Characteristic of Cu-based Metallic Glasses ……………………… 7

2.4.1General information of Cu-based BMGs ……………… 7

2.4.2 Background information of the chosen elements ……… 7

2.5 Applications of Bulk Metallic Glasses ………………………….... 8

3. Experiment

3.1 Fabrication of Bulk Metallic Glasses …………………………… 9

3.2 Sample Preparation………………………………………………… 9

3.3 Confirmation of Amorphous Structure …………………………… 10

3.3.1 X-Ray Diffractometer (XRD) …………………………… 10

3.3.2 Differential Scanning Calorimetry (DSC) ……………… 10

3.4 Compression Test………………………………………………… 11

3.5 Vickers Hardness Test…………………………………………… 12

3.6 Scanning Electronic Microscopy (SEM) ………………………… 13

4. Results and Discussion

4.1 Thermal Analysis with DSC ……………………………………… 14

4.2 X-Ray Diffraction of BMGs ……………………………………… 17

4.3 Vickers Hardness Test …………………………………………… 18

VII

VIII

4.4 Compression Test ………………………………………………… 19

4.5 Fracture Surface under SEM ……………………………………… 22

4.5.1 Plane view………………………………………………… 22

4.5.2 Side view ………………………………………………… 24

5. Conclusions

6. Future Works

6.1 Prediction of Glass Forming Ability……………………………… 28

6.2 Morphology Study ………………………………………………… 28

7. References ………………………………………………………………… 29

Appendix (I-III)

1. Introduction

Bulk metallic glasses (BMGs) are alloys that have a variety of features not perceived in

conventional metallic materials. They have the mechanical behaviour between metals and

glasses which give a low shrinkage during cooling, resistance to plastic deformation,

resistance to wear and corrosion, but much tougher and less brittle than oxide glasses and

ceramics[1]. However, monolithic BMGs usually show poor plasticity and no strain

hardening ability during deformation at room temperature due to highly localized shear

bands. These weaknesses extensively limit the range of possible applications [2]. Based

on this reason, the development of bulk metallic glasses (BMGs) with improved

mechanical and physical properties has been among the most active research topics on

BMGs recently [3-5].

The discovery of BMGs, the mostly multi-component alloys, depends mainly on trial-

and-error. In general, a complex composition with improved glass-forming ability (GFA)

and mechanical properties is made by adding specific elements to a compositionally

simple glass-forming alloy. For example,he addition of Al in binary Zr–Cu or Zr–Ni

alloys can produce Zr-based BMGs with high GFA [3, 4]; Zr-based BMG with

micrometer-sized Nb particles exhibits high fractured strength and enhanced plastic

strain .[6]

In this dissertation, I will conduct a series of experiment to study the mechanical

behaviors (structural study in fractured area, deformation behavior, shear strain,

compression testes, etc) of four different compositions of Cu-based BMGs. I expected

that difference in mechanical behaviors among these specimens can be observed since In,

Si and Sn have different atomic size and mechanical properties. Also alteration of the

glass forming ability of different glassy sample alloys may be achieved.

1

2. Literature Review

2.1 Background of Metallic Glasses

Metallic glass is known as amorphous metal which is a metallic material with a

disordered atomic-scale amorphous structure, like glasses. The first breakthrough in

metallic glass formation came in 1960 when Klement, Willens, and Duwez discovered

that Au75Si25 could be made amorphous by rapid cooling from liquid state [7], as depicted

in Fig 2.1; crystalline phases were bypassed through rapid cooling. Under rapid

quenching of the molten metal, nucleation and growth of crystals are kinetically forbade

and the molten metal retains in a configuration of frozen liquid. Since then, extensive

studies on metallic glass have been started. In 1974, Chen [8] first reported of BMGs

ternary Pd–Cu–Si metallic glasses with millimetre scale, cylindrical rods 1 to 3 mm

diameter, at a cooling rate of 103 K/s. In the early 1980s, Turnbull and co-workers [9, 10]

succeeded in forming the well-known Pd–Ni–P bulk metallic glasses by using boron

oxide flux method. In the late 1980s, Inoue et al. [11,12] successfully discovered new

strongly glass forming multi-component alloys containing mainly common metallic

elements with large undercooling and lower critical cooling rates. Since then, continuous

casting processes for commercial manufacture of metallic glasses ribbons, lines, and

sheets were developed [13]. BMGs with advanced properties, high glass formability and

were continuous published throughout the decades.

Fig 2.1 Isothermal formation of a metastable metallic glass by interdiffusion at

2

temperature T d between two crystalline phases A and B when the stable phases A3B and

AB do not nucleate. T g =glass transition temperature. [14]

2.2 Fabrication of Bulk Metallic Glasses

“Bulk” metallic can be defined as one which can be quenched from the melt into a

specimen with critical cooling rates low enough (approximately 102 K/s) to allow

formation of amorphous structure in thick layers (over 1 millimeter). They can be

characterized by some empirical rules:(a) Deep eutectics in phase diagrams of alloys, (b)

a large value of ΔT,(Tx-Tg), (c) multi-component alloy systems consisting of more than

three elements, and (d) a large different in atomic size.[15]

2.2.1 Critical cooling rate

Cooling rate is one of the most important factors for fabricating metallic glassy with fully

amorphous structure. Fast cooling rate can suppress the nucleation and growth reaction of

a crystalline phase in the supercooled liquid region and deviate from it , that is the

interval between the glass transition temperature (Tg) and crystallization temperature (Tx)

for the formation of amorphous alloy. Critical cooling rate (Rc) for glass formation can be

calculated by using Uhlmann [16] criterion for glass formation: an isothermal

transformation curve for a volume fraction of 10-6 is constructed and the rate is taken to

be undercooling at the nose of the curve divided by the transformation time at that

temperature. When a metallic glassy with low critical cooling rates have been discovered

in 1969 by H.S. Chen [17], a term Glassy Forming Ability, GFA, was introduced and

critical cooling rate (Rc) was used as a parameter to indicate the GFA. The minimum

critical cooling rate for Fe-, Co- and Ni-based metallic glass is only 102K/s for Pd-Ni-P

and Pt-Ni-P amorphous alloy [18]. Recently, it has been found that the critical cooling

rate can be as low as 0.1K/s for multi component amorphous alloy [19].

2.2.2 Composite atomic size

Metallic glasses exhibiting high glass forming ability can be considered as alloy phases

with specifized compositions. Negal and Tauc in the 1970s addressed the formation of

the metallic glasses that consist of noble and polyvalent metals by examining their

3

electronic structure [20]. Atomic size is also a significant factor in determining glass

forming compositions. Amand and Giessen, after examining alkaline earth systems,

pointed out that liquid viscosity as well as the amorphous alloy formation are influenced

by the difference of atomic sizes [21].

2.2.3 Suction casting mechanism

A cylindrical sample is set up by sucking the molten alloy into a copper mold through

suction force caused by the difference in gas pressure between melting chamber and

casting chamber, as shown in Figure 2.2. Immediately before casting, a piston with a

diameter of 16mm which was set at the center of the copper mold for arc melting is

moved at a high speed of 5.0 m/s and sucking force is generated by the rapid movement

of the piston.

The most commonly used preparation methods are copper mold casting and arc melting.

However, using the arc melting method, it is very difficult to completely suppress the

precipitation of the crystalline phase because of the ease of the heterogeneous nucleation

formed in different part of the copper mold [14].

Fig2.2. Simplified mechanism of arc melting suction casting method [14]

2.3 Mechanical Properties of Bulk Metallic Glass

Due to the random structure of amorphous alloys that lack crystal defects, bulk metallic

glass possess better mechanical properties than those of their crystalline alloys.Fig2.3

4

shows the different in atom arrangement of crystalline and amorphous structure upon

deformation.

2.3.1 Compression Test at room temperature

In uniaxial compression test, metallic glasses exhibit higher yield strength than crystalline

metals and fail in a brittle way with little or even no plasticity. Tensile strength of 2 GPa

[22] and toughness of up to 55MPa/m [23] have been reported. When they are stressed

beyond the yield strength, they tend to undergo localized shear due to lack of strain

hardening. Thus, even though they are not generally deformable like crystalline alloys,

they resist fracture.

2.3.2 Plastic deformation

At room temperature, BMGs tend to form shear bands with localized plastic deformation.

[24]. As these bands are favored sites for further plastic flow due to strain softening, they

normally lead to the failure that generally breaks a sample along a single shear band [25].

BMG-base composites have been developed by introducing ductile crystalline phases into

BMGs. With the crystalline phases, the composites can have dislocation related work

hardening behavior that can suppress the strain softening of a single shear band and

promote the generation of multiple shear bands in glassy matrixes. To enhance the

ductility of BMGs, several trials have been made to customize the microstructure by

strengthening the glassy structure. Several means have been invented such as introducting

nanocrystalline [26, 27] precipitates, e.g.Nb [27], shrink-fit copper sleeve [28] formed

through partial crystallization by exposing the as-prepared BMGs to an annealing

treatment.

The metallic glass does not have any defects that facilitate plastic flow in crystalline

materials, such as dislocations or gain boundaries. Without microstructural features to

direct and distribute the flow, severe shear bands associated with localized decrease in

glass viscosity form and propagate unhinderedly through the material. The possibility of

disastrous failure associated with the rapid propagation of shear bands is one of major

concerns when using BMGs in structural applications where reliability is critical. Both of

5

strain localization and shear band propagation are the exceptional problems under tensile

stress states because failure may occur along a single shear plane almost without

exhibiting any prior measurable plastic deformation.

Figure 2.4 shows the typical strengths and elastic limits for various materials [29].

Metallic glasses are typically much stronger than crystalline metal counterparts (by

factors of 2 or 3), are quite strong (even more so than ceramics), and have very high

strain limits for the elasticity. It can be seen that the elastic energy required to yielding

for bulk amorphous alloy is much larger than that of crystalline alloy.Again, the graph

shows metallic glasses posses a limited ductility, which is potentially a significant barrier

to their widespread application.

Fig2.3 Theoretical diagrams indicate different atomic arrangement of glassy and

crystalline structures upon shear force. [30]

6

Fig2.4 Typical strengths and elastic limits for various materials

2.4 Characteristic of Cu-based Metallic Glasses

2.4.1General information of Cu-based BMGs

Among different composition of BMGs, Cu- and Zr- based are the most common. This is

due to the high glass forming ability of them. Also, choosing appropriate compositions

can alter the minimum cooling rate of amorphous alloy and thus affect the GFA of certain

BMGs. This is why Cu- based MGs with different composites have been invented and

studied.

The recently developed Cu-based (more than 50% Cu content) BMG is reported to have

high tensile strength in the range of 2.0-2.8GPa. Also, there are some reports indicating

that the Cu-based bulk amorphous alloys usually possess high GFA [31] and good

ductility [32].

2.4.2 Background information of the chosen elements

The selection of additional elements is bases on two aspects. The first criterion is the

chosen elements should be within similar group, e.g. group V and VI. The second is to

consider the intrinsic mechanical properties which may give influence on the behavior of

final amorphous alloy, e.g. Silicon and Tin have distinctive mechanical properties.

7

Figure 2.5 illustrates schematically the variation on the atomic size of the three elements.

Their atomic radii are 159pm for Zr, 125pm for Al, 155pm for In, 145pm for Sn, 111pm

for Si and 128pm for Cu. Table 2.1 shows the fundamental information of the fourth

chosen elements in this study. One guiding principle of designing BMG alloys is to use

elements with significant differences in atomic size, which leads to a complicated

structure that retards the movement of atoms as well as the crystallization [33].

Crystalline materials with body-centered cubic structure, which generally have relatively

high hardness, may help to limit the movement of the shear band in effective manner.

The superior mechanical properties and its relatively low price make them suitable as

new engineering materials. [34]

Fig2.5 Schematical atomic sizes of the chosen elements

Electron

configuration

Density

(g/cm3)

Crystal structure Young's

modulus

(GPa)

Atomic

radius

(pm)

Silicon (Si) [Ne] 3s2 3p2 2.33 Diamond cubic 150 111

Tin (Sn) [Kr] 4d10 5s² 5p² 7.265 tetragonal 50 128

Indium (In) [Kr] 4d10 5s2 5p1 7.31 tetragonal 11 155

Table 2.1 Fundamental information of the chosen elements

2.5 Application of Bulk Metallic Glasses

Owing to the unique properties of bulk metallic glasses, they have many potential

applications. Among the most important ones are: superior strength and hardness,

shaping and forming in viscous state, exceptional corrosion resistance, reduced sliding

8

friction and enhanced wear resistance. These properties make it feasible for applications

in near-shape fabrication by injection molding and die casting, joining and bonding,

coatings, biomedical implants, and synthesis of nanocrystalline materials.[35] Beyond

opening the door for high strength applications, BMGs are excellent materials for

mechanical energy storage applications (e.g. springs) and act as soft magnetic materials

for common mode choke coils[30]. Also the development of bulk metallic glasses with

soft magnetic properties would be beneficial for energy conservation.

9

3. Experiment 3.1 Fabrication of (Cu46Zr46Al8)99X1

The bulk Cu-based metallic glasses were prepared by arc melt suction casting method. A

Parameters were set by addition of Si, Sn and In to the referencing Cu-base bulk metallic

glass, Cu46Zr46Al8. Series of (Cu46Zr46Al8)99X1 (where X= Si, Zn and In) and the

referencing alloys were prepared by arc melting the mixture of the respective pure

elements with purities higher than 99.9% in appropriate proportion under a Ti-gettered

argon atmosphere. Four Cu-Zr-Al-X compositions were designed which satisfy 99 to 1

atomic ratio and their weight compositions are listed in Table 3.1. Each ingot was re-

melted for four times in the arc melter under vacuum as to obtain chemical homogeneity.

Bulk alloy rods with 2 and 3 diameters and ~30mm length were prepared by suction

casting into a water-cooled copper mold in a purified argon atmosphere. An inert

atmosphere for casting can reduce oxygen contamination on the as-cast samples.

(Cu46Zr46Al8)99X1

Weight (g) Zr Cu Al X

X=Si 5.247 3.655 0.270 0.035

X=Sn 5.061 3.525 0.260 0.145

X=In 5.101 3.553 0.262 0.141

Table 3.1 Weight composition of different specimens

3.2 Sample Preparation

Total of 8 ingots in which 2 samples were fabricated within the same parameter. The

2.8mm bulk samples were cut into 3 portions, as shown in Fig.3.1 in Appendix II, by a

low speed diamond saw lubricated with oil. Sectioned samples were cleaned by acetone

and followed by ethanol in the ultra-sonic bath for 5 minutes. The samples were ground

flat to parallel with Silicon carbide paper from 240 down to 1200 grit size. For those

prepared for hardness test, samples were hot mounted with acrylic and further polished

with 5 μm aluminium oxide powder to mirror surface.

10

3.3 Confirmation of amorphous structure

Two different techniques were employed to confirm the microstructure of the as-cast

samples.

3.3.1 X-Ray Diffractometer (XRD)

The amorphous nature was identified by the Siemens D500 XRD with Cu Kα radiation

sources (λ=1.5406Å). Samples being tested were cut into 1 mm thickness by low speed

saw and washed with acetone for 5 seconds. The operation voltage and the current of the

x-ray were 30mA and 40KV. The detector was driven to scan from 2θ of 20 to 80 degree

with step size and time of 0.02°and 1 second respectively. A broad diffuse peak without

sharp point should be observed when amorphous structure is being tested. Fig 3.1 depicts

a typical XRD pattern of the referencing sample.

3.3.2 Differential Scanning Calorimetry (DSC)

The thermal properties of the as-cast samples were measured with a Perkin Elmer DSC 7

under a purified nitrogen atmosphere. Each sample rod with diameter of 3 mm was cut

into fine layer (thickness~1 mm) by a slow speed diamond saw and the samples were

weighted between 15 mg to 20 mg. The glass transition temperature (Tg) and

crystallization temperature (Tx) of the amorphous alloy were determined by the DSC 7

which was first calibrated with the melting transition of tin. The DSC scanning range was

made from 323K to 873k at a constant heating rate of 20 K/min. The samples are

expected to shows Tg and Tx within this range.

Metallic glasses exhibit distinct glass transition followed by super-cooled liquid region,

and then exothermic reaction in DSC curve. They show an endothermic heat event

characteristic of glass transition followed by a characteristic exothermic heat releases .

These indicate the transformation from the metastable undercooled liquid sate into the

crystalline compound. The glass transition temperature (Tg), onset crystallization

temperature (Tx) were marked in the DSC trace. [36, 37]

11

Fig 3.2 DSC curve of the Cu46Zr46Al8 at heating rate of 20K/min

3.4 Compression Test

Series of compression tests were performed using Zwick/Z030 tensile tester at room

temperature, as shown in Appendix II. The strain rate was set to 0.03mm/min and the

stress-strain curve was obtained. All the samples were clamped by a flat stainless steel

mount and grounded with silicon carbide paper to ensure that the upper and lower

surfaces of the sample were flat.

Sun et al. report that SEM observation of BMG casted reveals that there exists an obvious

microstructure transition from the center region (existence of nano-crystal) to the edge

(fully amorphous structure) due to the different cooling rate, In order to verify the effect

of cooling rate inside the copper suction tube, the as-cast samples were cut into 3

different parts, namely Top, Middle and End, with L/D ratio of 2, as shown in Figure 3.3.

12

Figure 3.3 Three Portions of upper 2mm and lower 3mm diameter fabricated bulk

metallic glasses.

The fracture surface morphologies of the deformed specimens will be observed by SEM.

The compressive strength, calculated elastic modulus and compressive strain will be

evaluated from the stress-strain graph. In Sn and In containing Cu-Zr-base alloy, it is

expected to see if there is any improvement in the plastic deformation and increase in

hardness for Si-BMG.

3.5 Vickers Hardness Test

Vickers hardness test uses a diamond pyramid indenter, which is pressed into the surface

to be tested using a prescribed force load for a specific amount of time. After the indenter

has been removed, the size of the indentation left behind is measured and calculated as

the hardness of the material. Soft materials will show large indentation; hard materials

display small indentations.

Samples were mounted with 3mm diameter acrylic and further polished with 5 μm

aluminium oxide powder to mirror surface. A Vickers Hardness Tester FV-700 was

employed at loads of 1kgf, 10kgf and 30 kgf with dwell time 10s. Each sample was tested

and the effects of changes in indentation load on the hardness and appearance of indents

13

was investigate. The Vikers hardness test of 1kgf, 10kgf and 30 kgf was taken from 10, 3

and 3 random selected points respectively and the average value was used.

3.6 Scanning Electronic Microscopy SEM

The SEM measurement was performed using a SEM-JEOL JSM-820 scanning

microscope operated at 20kV. Both plane view and side view of the fracture surfaces

were observed.

The fracture samples with a shear angle were washed with alcohol in order to remove any

impurities and dusts on the fracture surface. The magnifications of SEM ranges from

x300 to x1500.

14

4. Results and Discussion

4.1 Thermal Analysis with DSC

The most distinctive physical property that defines a glassy solid is a glass transition. The

glass transition temperature Tg is defined as the onset temperature of the endothermic

event. It can be detected by the rapid increase of the heat capacity over a small range of

temperature [38]. To determine the values of Tg from the DSC graph, the curves were

enlarged along the heat flow coordinate to clarify the flections due to glass transition. The

crystallization temperature Tx is defined as the onset temperature of the exothermic event.

e.g. the Tg and Tx of Cu46Zr46Al8 are 458.9K and 528.9K respectively. The accuracy of

DSC temperature measurement is about 1 K and these curves give confirmation of

amorphous structure to the ingots.

Figure 4.1 summarize the DSC traces of the reference Cu46Zr46Al8 BMGs and those with

1at % Sn, In, and Si at heating rate of 20K/min. The traces show structure characteristic

change of endothermic glass transition and then followed by a supercooled liquid region

( T△ = Tx – Tg), an exothermic crystallization peak due to crystallization and an

endothermic reaction due to melting is ensued. Single exothermic peak is observed in

every curve indicates that crystallisation process occurs once only. This gives the

evidence on the single phase of the amorphous structure. From the result obtained, there

is a large supercooled liquid region for the referencing Cu-Zr-Al BMG, implies a stability

of super-cooled liquid state. The Tg of referencing BMG Cu46Zr46Al8 is the highest

among the four BMGs, which is 458.9K. The lowest is the one with 1 at% of indium,

which is 451.5K. A summary of the thermal analysis data deduced from these curves is

tabulated in table 4.1.

15

Fig4.1 DSC curves of the reference Cu46Zr46Al8 and with 1at % Sn, In, and Si range from

400K to 580K

Tg (K) Tx (K) ΔT(K)

Cu46Zr46Al8 458.9 528.9 70

(Cu46Zr46Al8)99Si1 452 531.4 79.4

(Cu46Zr46Al8)99Sn1 457.8 515 57.2`

(Cu46Zr46Al8)99In1 451.5 511.5 59.5

Table 4.1 Thermal parameters of (Cu46Zr46Al8)99X1 metallic glasses

The analysis of the data shows: 1. Up to ~440K the alloy structure remained amorphous;

2. an endothermic knob at ~440-460K is related to glass transition; 3. All Tg supressed

by the addition of the fourth elements. 4. A transformation (crystallization) from

amorphous to crystalline state ranges from ~510-530 K.

16

From the DSC summary in Table 4.1, with the addition of 1% at Sn and In, both Tx and

Tg decrease within a range of ~8K and the ΔT decreases from ~70K to ~60K. The

degradation of ΔT implies that with addition of In and Sn, the thermal stabilities are all

reduced. These can be explained by the small atomic size different with addition of Sn

and In.

In contrary, an addition of 1 at% of silicon shows a reverse effect when compared with

the referencing Cu46Zr46Al8BMG., in which the Tg was suppressed by 6K and Tx

increased by 2.5K, therefore ΔT increase by ~9K, This result is consistent with the

increased stability of the undercooled liquids by additional of 1 at% Si to Cu-Zr-Ti BMG

[39] and Boron additions increases thermal stability of Zr-Cu-Al alloy system.[40]

On the basis of the empirical BMG formation criteria, the formation of bulk metallic

glasses should satisfy multi-component alloy systems and large differences in atomic size

between constituent elements. Since the addition of Si give a significant increase in

atomic size difference, compared to Sn and In constituent elements. This is in agreement

with the above two rules for the empirical BMG formation criteria. According to the

argument for the size effect, composition of atoms in larger atomic size difference can

enhance the thermal stability. The thermal stability predicted by the difference in atomic

size show there is an increase in ΔT by the addition of Si. Therefore, the prediction of

GFA by applying atomic size difference theory is appropriated.

Besides, it is well known that crystallization is controlled by diffusion of atoms. Because

of the highly dense arranged and complex microstructure of atoms, it causes slow atomic

mobility and long-range atomic diffusion; an increase in the dense random packing of the

supercooled liquid, making atomic diffusion more difficult [41]. This explanation also

proved to be valid by the experimental result.

17

4.2 X-Ray Different (XRD)

Figure 4.2 shows the X-ray diffraction data of the as-casted samples in the range 2θ of

20° to 80°. The broad diffraction maxima between 35° and 45° reveal the amorphous

structure of the sample. The pattern shown underneath well illustrated the ability to

obtain glassy structure with the addition of the fourth element (Si, Sn, In) is not

suppressed. Also, the absent of sharp crystalline peak in the diffraction pattern gives

evidence on the glassy state of the alloy. However, due to the limitation of XRD, small

volume factions of crystalline phases in the amorphous matrix may not be detected.

Fig 4.2 XRD patterns of Cu46Zr46Al8, (Cu46Zr46Al8)99Si1, (Cu46Zr46Al8)99Sn1 and

(Cu46Zr46Al8)99In1 alloys.

18

4.3 Vickers Hardness Test

The Vickers hardness value (Hv) is one of indicators to evaluate the resistance to plastic

deformation. Ductility can be evaluated by the mean of standard hardness tester for small

brittle sample. Three different loading; 1kgf, 10kgf and 30 kgf were employed purposely

to observe any collapse or crack traces near the indentation surfaces. Enbrittlement of the

specimens can be evidenced by the occurrence collapse at the edge of indentation.

The average dependence of the Vickers hardness (Hv) of the samples is summarized on

Table 4.2. Hardness of the tested samples is in the order of: (Cu46Zr46Al8)99Si1>

(Cu46Zr46Al8)99Sn1> Cu46Zr46Al8> (Cu46Zr46Al8)99In1. Among these alloys, collapse edges

can be observed only in the (Cu46Zr46Al8)99Sn1 alloy. It is well-known that crystalline

alloy is more brittle than amorphous alloy; the ductile to brittle behavior of BMGs can be

shown by the appearance of indentations. From the Fig 4.3a, it is obvious that collapsed

edge shape appeared on the surface near the pyramidal indentation point. 7 out of 10

indentations exhibit collapses at 1kgf load of (Cu46Zr46Al8)99Sn1. Therefore, it is

suspected that the existence of nano-crystals in the amorphous matrix. However, no

collapse due to the embrittlement can be observed for 10kgf and 30 kgf.

Fig 4.3 (a),(b) Optical microscope photograph of 1kgf indentation of Cu-base BMG with

1 at% Sn show the appearance of collapsed edges (Pointed by arrow).

19

For the crystalline alloys, hardness is directly related to the yield strength of a material.

The yield strength of a crystalline alloy is determined by dislocation interaction with

lattice defects such as free volume, dislocation and grain boundary. For amorphous alloys,

due to the lack of crystalline defects, the hardness is relative low compared with

crystalline alloys. In addition, the properties of fundamental elements of the alloy also

contribute to the mechanical properties of the alloy. Silicon is considerably harder than

Indium and Tin. The result of increase in hardness by addition of Si might be attributed to

the hard nature of silicon. However, addition of Sn does not give significant reduction in

hardness of the alloys. Explanation on the hardening effect of Sn addition by other

approach has to be employed, e.g. existence of nano-crystal phases.

4.4 Compression Test

The compressive mechanical properties at room temperature of the suction cast Cu-base

BMGs have been measured by using Zwick/Z030 tensile tester, shown in Appendix II, at

room temperature. 3 portions of each ingot are cut with height to diameter rations of 2:1

and conducted uni axial compression test. The strain rate of compression tests is set at

0.03mm/min.

From the Figure 4.4(b)-(e) stress-strain curve shown in Appendix I, referencing

Cu46Zr46Al8, (Cu46Zr46Al8)99Si1 and (Cu46Zr46Al8)99In1 glassy alloys do not show a distinct

plastic deformation behavior except for the top part of (Cu46Zr46Al8)99Sn1 ingot. The

yield strength, young’s modulus and plastic strains of all samples are summarized in table

4.4. Among these samples, (Cu46Zr46Al8)99Si1BMG gives the highest compressive

strength which is 2.07 GPa followed by addition of In, Sn and the lowest is the

referencing BMG, which are 1.93GPa, 1.88GPa and 1.74GPa respectively. Both the

sample exhibited a liner elastic elongation within 5% and only Sn-BMG shows a trend

followed by a plastic elongation of about 8.1%, other samples fracture immediately at

yield stress. The referencing Cu46Zr46Al8 BMG shows no evidence of macroscopic

yielding and plasticity, as reported for the most Cu-base BMGs.

20

Young’s modulus gives evidence on the stiffness of the testing samples. It is defined as

the ratio, for small strains, of the rate of change of stress with strain. This can be

determined from the slope of a stress-strain curve created during compression tests. From

the stress-strain graphs obtained, addition of silicon increases the young modulus of the

Cu-Zr-Al BMG, from 23.03 to 37.55GPa. This phenomenon can be explained by the stiff

nature of silicon element. Moreover, BMGs with addition of In which have a soft nature

show an decrease in stiffness compare with the referencing BGM. These results are

consistent with the hardness test which give the stiffness of the sample with the strongest

(Cu46Zr46Al8)99Si1> (Cu46Zr46Al8)99Sn1> Cu46Zr46Al8> (Cu46Zr46Al8)99In1. However,

addition of Sn increase the young’s modulus by ~2GPa compare with the referencing

BMG, this cannot be explained by the nature of the addition element and may due to the

complicated amorphous structure of the glassy alloy formed.

In order to further explain the hardening effecting by addition of silicon, model lead to

strong softening in the shear band proposed by other researcher should be mentioned.

Argon [42] introduced the concept of shear transformation zone. These zones begin as

small region where the local atomic structure is capable of rearrangement under a given

applied shear stress and the ability of a region to undergo a shear transformation depends

in the local atomic density. This theory is valid for the addition of small atomic size

silicon which fills in the free volume of the amorphous matrix and forbids the initiation

of atomic dislocation upon compression. This gives raise to the higher yield strength of

Cu-Zr-Al-Si alloy system. However, the atomic sizes of the fourth element Sn, In are of

similar size with the rest of the referencing system, no significant increase in strength can

be observed.

Figure 4.4(a) shows the true stress-strain curve of top part of (Cu46Zr46Al8)99Sn1 ingot.

The compressive yield stress, elastic modulus and plastic strain are 1.88 GPa, 25.49GPa

and 8.1%, respectively. The upper part on the Cu-Zr-Al-Sn rod exhibits a large plastic

deformation and strong increase in strain, compared with other alloys in this study. After

yielding, the Sn-BMG shows stress increase with strain indicating a work hardening up to

8.1 %. It is suspected that the atomic-scale heterogeneities exhibited by the structure of

21

the metallic glass facilitate nucleation and continuous multiplication of shear bands. This

can be explained by the interaction and intersection of shear bands increases the flow

stress of the material that causes further deformation, displaying a ‘work hardening’-like

behavior.[43] Therefore, an investigation on the shear band evolution during deformation

will give conformation to the ductility of (Cu46Zr46Al8)99Sn1.

Focus should be put on the Cu-Zr-Al-Sn alloy system, since only the top part of the ingot

exhibit a distinct plastic deformation of 8.1 %. One of the possible reasons for this

phenomenon is the discrepancy of cooling rate inside the Copper mold. The upper of the

ingot is very close to the arc and did not go through the whole suction tube; therefore,

lower quenching rate and relatively higher temperature resulted at the upper part. This

gives raise to the formation of XRD undetectable nano-crystals imbedded in the

amorphous matrix. Further confirmation of the crystal structure by TEM should be

carried.

Average Yield Strength

(GPa)

Young’s modulus

(GPa)

Plastic strain

(%)

Hardness

(Hv)

Cu46Zr46Al8 1.74 23.03 - 531.62

(Cu46Zr46Al8)99Si1 2.07 37.55 - 562.85

(Cu46Zr46Al8)99Sn1 1.88 25.49 8.1 553.78

(Cu46Zr46Al8)99In1 1.93 20.88 - 506.48

Table 4.2 Summarize the calculated yield strength, young’s modulus, plastic strain and

hardness of all samples.

22

Fig 4.4(a) True stress–strain curves of upper part of (Cu46Zr46Al8)99Sn1 ingot under

uniaxial compression. Inset shows nominal stress–strain curves of (Cu46Zr46Al8)99Sn1with

three portions of the ingots.

4.5 Fracture Surface under SEM

The plane and side views of fractured surfaces were observed under SEM operated at

20kV.

4.5.1 Plane-View

Figure 4.5 (a)-(b) and (c) from Appendix III show the morphology of the fracture surface

of (Cu46Zr46Al8)99Sn1 by SEM. From the result, a clear and well-developed vein-like

pattern spread over the fracture surface following the shear direction. The molten vein-

23

like pattern, evidence of a liquid–like layer where the glass has softened significantly, is a

region with a considerable high temperature (>900 K) due to local heating. In addition,

the remarkable development of the vein-like pattern and the significant increase in the

diameter of the melt veins suggest that the temperature during the final fracture increases

because of the suppression of the final fracture resulting from the good ductility. [44]

Furthermore, the increase of the diameter of the veins also implies the increase of the

thickness of the shear deformation region which causes an increase of the energy required

for plastic deformation and final fracture [45].

Differences were noted in the fractured surface appearance between the referencing and

1at% Sn BMG. Higher dense and well developed molten vein pattern can be found in

referencing Cu46Zr46Al8 suggesting the heat dump from the stored elastic energy could be

contributing to fast crack propagation by allowing the crack to propagate along a hotter or

even molten path.

(a) (b)

Fig 4.5 (a) (b) SEM images of the vein-like pattern of (Cu46Zr46Al8)99Sn1

(b) Shows a site with higher dense of vein pattern at the core.

24

4.5.2 Side View

On the lateral surface of the fractured specimens, shear bands (SBs) with different

orientations can be observed. In general, all fracture rod show few primary and secondary

shear bands expect for the one with 1 at% Sn, which have a plasticity of 8%, shows

clearly high dense SBs along the cross section as shown in Figure 4.7 (a)

As shown in Fig 4.6, numbers of large localized SBs together with high dense secondary

SBs were formed during plastic deformation, therefore the sample did not result in rapid

fracture like other glassy alloys in this study. From the SEM image, many narrow

secondary SBs have been triggered near the large primary SBs and some of them are

intersected with other primary SB, leading to a considerable plastic strain without

catastrophic failure. Many secondary SBs are kinked, bifurcated, branched and interacted

with others or with the primary SB, again, indicating the resistances exist to hinder the

rapid propagation of shear bands. Moreover, smaller inter-shear bands space and more

branched shear bands suggesting SB multiplication yield a global ductility. [46]

The activation of secondary SBs would reduce the local stress concentration and reduce

the propagating rate of the primary SB and thus, reduce the possibility of rapid fracture. It

also implies that the Cu-Zr-Al-Sn alloy has a strong resistance to crack nucleation and

propagation. Furthermore, slipping occurs in some of the SBs with a shear offset of

5.7µm shown in the inset of Figure 4.6, this phenomenon also contribute to the plastic

deformation of the Cu-Zr-Al-Sn sample. Ref [46] shows similar result in Ti-Cu-base

metallic glasses indicating the existence of nano-particles in glassy matrix trough TEM.

Serrated flow of shear band due to viscosity can also be observed on 1at% Sn specimen

and these greatly increase the resistance of metallic glass to deformation. There has been

considerable debate in the literature as to the cause of this ductility. Two hypotheses

emerged to explain the local changes in viscosity occur in shear bands: free volume

generation and localize adiabatic heating [45, 46]

25

Comparing the one with 1at% Si addition which shows low ductility, localized shear

bands were formed with fewer secondary SBs intersecting, crack nucleate and propagate

mainly through the dominant primary SBs, the specimen thus catastrophic fracture

rapidly without plastic deformation and strain softening, as shown in Figure 4.7 (a) (b). In

contrary, if the crack nucleation or propagation is hindered, fracture will be hindered too,

resulting in higher ductility. [47]

Das [48] and Liu [49] both found that introducing inhomogeneous amorphous phases can

mostly improve the plasticity in the bulk metallic glasses. Lee [50] also considered that

the embedded crystalline granules in a Cu-based BMG can enhance the plasticity

extensively. These results show that two coexisting phases can influence shear band

movement and enhance the plasticity or work-hardening of BMGs.

5.7µm

Fig 4.6 SEM images shows high dense shear bands with primary and secondary shear

bands. Shear offset of 5.7µm shown in the inset image of sample Cu-Zr-Al-Sn alloy.

26

Secondary shear band

Primary shear band

(a)

(b)

Fig4.7 SEM images of shear bands formed during uniaxial compression of

(a) (Cu46Zr46Al8)99Si1, (b) Cu46Zr46Al8. Both with arrows indicate the direction of

secondary SBs propagation.

27

5. Conclusion In this study, the successful fabrication of 2 and 3 mm diameter amorphous alloys of

(Cu46Zr46Al8)99In1 and (Cu46Zr46Al8)99Si1 is proved by the peaks in DSC curves and broad

bend patterns in XRD curves. The enhancement of thermal stability of Cu-Zr-BMG with

1% atomic silicon addition is demonstrated by the enlargement of T in DSC trace.△

From the evidences in stress-strain curves and morphologic studies of the fracture

surfaces, the addition of Sn plays a great important role on the improvement of ductility

for the Cu-Zr-Al amorphous alloying system. The plasticity of BMGs can be extensively

improved by 8.1%. It is suspicious that nano-crystals exist in the upper part of

(Cu46Zr46Al8)99Sn1ingot. It again, proves the cooling rate is critical for the formation of

fully amorphous alloys.

In the suction tube, slightly difference in quenching rate lead to the formation of nano

crystals on the top of the ingot. It implies that it may facilitate the precipitation of

nanocrystallites in the amorphous phase by introducing a small amount of Sn when the

BMGs are deformed under compression [51]. As a result, under SEM examination,

propagation of the primary shear bands can be impeded by nano-crystals and new shear

bands have to be initiated. The formation of high dense of shear bands thus gives

explanation on the large plasticity performance of Tin containing glassy alloy.

Meanwhile, working hardening was shown in the Vickers hardness test with the addition

of silicon element. This can be explained by the hard nature of silicon atom. Collapsed

indented edges observed in Sn- containing BMG illustrate the embrittelemnt effect of 1%

atomic addition of Tin.

28

6. Future Works

6.1 Prediction of Glass Forming Ability

The melting behaviour of the samples can be investigated by differential thermal analysis

(DTA). Obtaining traces with the solidius temperature Tm and liquidus temperature Tl

which are respectively the onset and end temperature of the endothermic melting, is a

more precise means to estimate the glass forming ability of BMGs. The calculated values,

i.e., Trg(=Tg/Tl), γ[=Tx/(Tl + Tg)] and δ [=Tx/(Tl −Tg)]and the values of the maximum

glassy sample diameters, D max can be used together as indicators. The higher the values

of Trg or δ, the larger the maximum glassy sample diameters, and consequently, the

higher is the GFA.

6.2 Morphology Study

In this study, the crystallization behaviors are examined by XRD and DSC. However, due

to the sensitivity of XRD and heat range limitation of DSC, nano-size crystal may not be

detected. in order to obtain more information about the bulk amorphous alloys, the

Dynamic Mechanical Analysis (DMA) should be considered since it is another effective

method to quantify the the viscoelastic nature of metallic glass due to plastic flow, so that

structural relaxation behavior of amorphous alloy can be investigated in more depth.

Besides, structural investigation by Transmission Electron Microscopy (TEM) should be

employed in the prospect to obtain detail in the microstructure, morphology and grain

size distribution. In the XRD results, the unidentified crystalline phase can be confirmed

by the Selected Area Electron Diffraction (SAED) technique.

29

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Appendix I

Figure 4.4 Nominal Stress-Strain cures for the different samples. (b) Cu46Zr46Al8 (c)

(Cu46Zr46Al8)99In1 (d) (Cu46Zr46Al8)99Si1 (e) (Cu46Zr46Al8)99Sn1

Appendix II

Figure 3.1 Three portions: top, middle and end from left to right were cut by diamond

saw

Fig 4.8 Vickers Hardness Tester FV-700

Appendix III

Fig4.5 (c) Plane view SEM image of Cu46Zr46Al8 fracture surface shows molten well-

developed vein-pattern. Arrow shows the direction of crack propagation. (c)

(d) (e)

Fig4.5 (d), (e) Side view SEM image of (Cu46Zr46Al8)99Sn1 fracture surface show molten

vein-pattern