COMPARATIVE STUDY Bi-lnSn TERNARY - TSpace … · COMPARATIVE STUDY OF Bi-InSn ... 2mm diameter by the Ohno Continuous Casting (OCC ... between the mold wall and the cast strand creates

COMPARATIVE STUDY OF Bi-lnSn TERNARY

EUTECTIC CAST WIRES

Surajit Sengupta

A thesis submitteâ in conformity with the requirements

for the deg ree of Master of Applied Science

Graduate Department of Metallurgy and Materials Science

University of Toronto

Q Copyright by Surajit Sengupta, 1998

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Dedicated to the memory of my mother

and uncle (Jethu)

COMPARATIVE STUDY OF Bi-InSn TERNARY

EUTECTIC CAST WlRES

BY Surajit Sengupta

Master of Applied Science

Graduate Department of Metallurgy and Materials Science

University of Toronto

1 998

Abstract A ternary eutectic Bi-lnSn alloy, which is brittle in nature and diffiwlt to

form by conventional methods, was successfully produced in the form of wire

2mm diameter by the Ohno Continuous Casting (OCC) process. The wire

produced has several unique features for example, superior surface finish, fine

microstruckire, uniform distribution of phases and consistent chemical

composition. In contrast statically cast samples exhibited segregation of bismuth

and double binary structures consisting of Bi ln~Sn and BiySn. In OCC samples

the matrix had a higher bismuth content and there was no evidence of

segregation. As a consequence of microstructural differences, wire produced by

the OCC technique had improved mechanical properties in terms of higher

strength and ductility compared to statically cast samples.

Acknowledgements

I would like to express sincere gratitude to my supervisors. Professor A. McLean

and Dr. H. Soda for their advice, encouragement and support throughout the course of

this thesis.

I also wouM like to acknowledge Dr. 2. Wang and Dr. J.W. Rutter for participation

in useful discussion.

I am grateful to Professor A. Ohno for the award of an Ohno Graduate Fellowship.

I am grateful for the financial support I received by University of Toronto Open

Fellowship and ALCAN award.

The technical support of Mr. F. Neub and Mr. Sal Boccia and the administrative

effort of the office staffs are deeply appreciated.

Finally I will ever remember the encouragement and the support from my wife,

Arpita; my son, Saswata and Babu-Ma.

Table of Contents

Page #

- Il

- 111

ABSTRACT - O

ACKNOWLEDGEMENTS

LIST OF FIGURES

LIST OF TABLES

CHAPTER 1: INTRODUCTION -

BASIC PRINCIPLE OF OHNO CONTINUOUS

CASTING - - O - CHAPTER 2:

CHAPTER 3: LITERATURE SURVEY - O O - 3.1 Lead alloys - O - O

3.2 Effect of bismuth on mechanical properties - 3.3 Microstructure of eutectic alloys - O

3.4 Segregation - -

CHAPTER 4: EXPERIMENTAL ASPECTS - O

4.1 Alloy preparation O O - Equipment O - O

Experimental procedure O

Static casting facility O - Experimental procedure for static casting

Sample grip design for mechanical testing

for tensile test

for microstructural

Sample preparation

Sample preparation

IV

observation-

CHAPTER 5:

CHAPTER 6:

RESULTS AND DISCUSSION - - O

5.1 Evaluation of casting condition for OCC - 5.1.1 The occurrence of breakout - - 5.1.2 Surface appearance of cast wire -

5.2 Microstructure - O - O

5.2.1 Microstructure of static cast rod - 5.2.2 Microstructure of OCC wire - -

5.3 Compositional uniformity in OCC samples - 5.4 Mechanical properties - O - - 5.5 Fracture surface of OCC and statically

cast samples - - -

CONCLUSIONS AND FUTURE WORK

6.1 Conclusions - O

6.2 Future work - - -

Page #

- 41

- 41

- 41

- 43

REFERENCES

LIST OF FIGURES Page #

Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Figure 6

Figure 7

Figure 8

Figure 9

Schematic diagram showing the principle of OCC and the

difference with conventional continuous casting - O

Different morphologies of binary eutectic alloys - - Different morphologies of ternary eutectic alloys - - Liquidus projection of Bi-ln-Sn system showing 350.5 K

ternary eutectic (24) - O - O

Schematic diagram of OCC equipment for generation of net

shape wires - O - O - Temperature profile during casting - - - O

Schematic diagram of resistance furnace for static casting in

a glass mold - - - - - - - Temperature profile inside the glass tube showing a plateau

region and a temperature gradient region - - O

Photograph of graphite split mold - O - Figure 10 Schemaüc diagram of resistance furnace for static casting in

graphite mold - - - O .. Figure 11 Temperature profile inside the mold cavity - - Figure 12 Photograph of INSTRON machine, mode18501 - Figure 13 Schematic diagram showing the problem associated with

INSTRON grip - - O O - - Figure 14 Photograph of modified grip, longitudinal and transverse

views - - - - 0

Figure 15 Schematic diagram of modified grip - - - Figure 16 Photograph of OCC wire and statically cast produds showing

significant difference in surface quality O

Page #

Figure 17 Cooling curve of the melt solidified inside the glass tube

showing the freezing point of eutectic alloy - 46

Figure 18 Backscattered SEM image of sample from the plateau region

of cast product in glass tube showing the phenornenon of

segregated blocks and complex regular structure of bismuth

(white phase) O - - - 47

Figure 19 Backscattered SEM image of sample from the temperature

gradient reg ion in glass mold showing directional solidification

and less segregation - O O - O - - 48

Figure 20 Energy dispersive X-ray (EDX) spectra showing,

a) Segregated bismuth,

b) Gray phase composed of bismuth and indium inside mottled

reg ion of bismuth complex regular structure,

C) Black dendrite composed of tin with a trace of bismuth,

d) Gray spine phase inside eutectic cell cornposed of bismuth.

indium and tin O O - 9 9 - 49

Figure 21 Backscattered SEM image of eutectic cell showing segregated

white bismuth phase in the fom of complex regular structure

around grain boundary, bismuth-indium gray spine and mottled

region of tin dendrite - - O O - 55

Figure 22 EDX analysis showing the presence of three elements

Bi, Sn and In in the decomposed structure of Bi complex regular

structure - - - - 9 - - - - 56

Figure 23 Backscattered SEM image showing a) the cubic, b) the fish

spine and c) the trigonal shaped complex regular structure of

bismuth - - - - - - 57

Page #

Figure 24 Backscattered SEM image showing the precipitation of

bismuth and tin in gray matrix and absence of gray region

where black tin phase is narrow - - - - Figure 25 Backscattered SEM image of complex regular structure - Figure 26 Backscattered SEM image of complex regular structure

Figure 27 Decomposed Sn dendrite showing several colonies or cells

having lamellar structure - - - - O

Figure 28 Backscattered SEM image of OCC wire, casting speed

14mm/min O O - - O - O


79mmimin O - O - - - O

Figure 30 EDX spectra of OCC wire showing,

a) White pure bismuth phase, b) Black dendritic tin phase,

c) Gray matrix phase without tin O - O


14mmiminl showing a) tin dendrite and associated bismuth and

b) at higher magnification O - - O - - 68

Figure 32 Plot to confirm compositional uniformity measured at different

Locations - - O - O - O - 72

Figure 33 Backscattered SEM image of statically cast rod accepted for

tensile test. Segregated bismuth was removed by machining

during sarnple preparation - - - - - - 76

Figure 34 Schematic diagram of tensile specimen - - - 76

Figure 35 Photograph of turnings to show difterence in material proopeities

between OCC and statically cast samples - - - 77

Figure 36 Comparison of elongation values between OCC and

statically cast samples at a crosshead speed of 1.25mmlmin

Figure 37 Cornparison of elongation values between OCC and

statically cast samples at a crosshead speed of 2.5rnrnfmin

Figure 38 Comparison of elongation values between OCC and

statically cast samples at a crosshead speed of 5mmlrnin


statically cast sarnples at a crosshead speed of 7.5mmlmin


statically cast sarnples at a crosshead speed of 10mmlmin

Page #

- 79

- 80

- 81

- 82

- 83

Figure 41 Plot of yield stress vs. crosshead speed showing the low value

of yield stress at higher crosshead speeds which produces

premature failure of the statically cast samples and the high

and consistent yield stress values of OCC samples - - 88

Figure 42 Plot of ultimate tensile stress vs. crosshead speed showing,

a) significant inconsistency in UTS for statically cast samples

at higher crosshead speed and b) higher and more consistent

UTS of OCC samples - - - - - - - 89

Figure 43 Plot of elongation vs. crosshead speed showing the greater

ductility of OCC samples compared to statically cast samples - 90

Figure 44 Fractography after tensile test at a crosshead speed of

1.25mmlmin. Both OCC and statically cast samples are ductile

in nature showing high reduction in area - - - - 93


2.5mmlmin showing no significant difference in reduction

of cross-sectional area - - Figure 46 Fractography after tensile test at a crosshead speed of

Smmlmin showing the evidence of ductility for both OCC

and statically cast samples - - O - - Figure 47 Fractography after tensile test at a crosshead speed of

7.5 mmlmin showing the ductile fracture of OCC sample

and brittle fracture of statically cast sample


Page #

- 94

- 95

- 96

l0mmlmin showing, a) high reduction in cross-sectional area

and ductile nature of fracture surface of OCC sample,

b) cleavages and no reduction in area on fracture surface

of statkally cast sample - - - - 97

Figure 49 Fractography of tensile test sample at a crosshead speed of

1.25mmhin and at high magnification - O - - 98

Figure 50 Fractography of tensile test samples, at a crosshead speed of

i0mmlmin and at high magnification - - - 99

Figure 51 SEM secondary image showing voids along the grain boundary

of a statically cast sample - - - - 100

LIST OF TABLES

Page #

Table 1

Table 2

Table 3

Table 4

Table 5

Table 6

Mathematical expression to obtain the dimension of the groove of modified grip according to diameter of the sample - 36

Optimum casting conditions for BI-ln-Sn alloy - -

Details of chemical composition of OCC samples -

Relation between crosshead speed and initial strain rate

Mechanical properties of statically cast samples -

Mechanical properties of OCC sarnples - - -

CHAPTER 1

INTRODUCTION

Ohno Continuous Casting (OCC) was developed at the Chiba lnstitute of

Technology in Japan. The name of this special type of casting was after the

pioneer of this process, A. Ohno. The idea of this process was to manufacture

alloys that are difficult to produce or that cannot be rolled, drawn, or extruded. In

conventional continuous casting the mold is cooled to solidify liquid metal inside

the mold to avoid run out of liquid metal at the mold exit. But the frictional force

between the mold wall and the cast strand creates surface defects. Apart from

surface defects the cast product processed through a cooled mold may have

different types of cast defects like shrin kage cavities, blowholes and surface

defects. So cast products have to undergo subsequent processing like surface

grinding for a better surface finish, annealing for microstructure and property

improvernent, hot or cold rolling for final required shape and desired properties.

The addition of further processing increases the cost of production.

There is a need for net shape cast products where the cross-sectional

area is very small and the high quality surface finish without any cast defect is an

advantage. There are different methods, the most common is wire drawing to

produce products at fast rate in wire form. But if the metal or alloys are brittle and

strain sensitive the drawing speed has to decrease drastically and sometimes it

cannot be produced. Soda et al (1) found that casting of bismuth which is fragile

and brittle in nature can be cast in wire fom by the OCC process. It should be

noted that bismuth expands during solidification and this would increase frictional

force during conventional casting. Wth the OCC proœss, single crystal bismuth

wire was produced which was ductile in nature. In other work Soda et al (2)

produced wire of diameter 1.7-2 mm with alloys having composition aluminum

1.5-7 wt % yttrium. The main feature of the wire was that it was unidirectionally

solidified with cellular or dendritic microstructure having constant unifonn

I

chemical composition along the length with excellent dimensional stability.

Considering the brittleness and small cross-sectional area, some aluminum

based alloys are difficult to produce or cannot be produced despite the fad they

are useful for surface hardening for the improvement in mechanical properties.

Again OCC was the only way to produce successfully aluminum with 25.50% Cu

and Al-Cu-Si alloys which are brittle in nature and ditficult to process in one-step.

It can be seen from the above brief discussion that with the OCC method

we can produce alloys which are brittle in nature and are difficult and in sorne

cases impossible to produce by conventional methods.

Several works have emphasized the need to avoid the use of lead in

solder alloys since it is toxic and will cause severe environmental problerns.

Lead-tin solder alloys are widely used since they are cheap with some

advantageous properties. The increasing use of electronic devices e.g. video,

audio, computer, telephone and wireless devices will cause lead pollution in

landfills, watennrays and soil. The recycling of the lead-bearing component may

not be economical and it is therefore important to investigate different lead-free

eutectic alloys (3). J. Glazer (4) has emphasized the urgency for development of

lead and cadmium free solder alloys. In this review it was mentioned that for the

42Sn-58Bi eutectic alloy at eutectic temperature, bismuth has significant

solubility (approximately 2iwt %) in ün. As a result, bismuth in the pure form

precipitates in the tin phase after solidification. Also, bismuth expands 3.87

volume % after solidification and alloy expands during solidification if bismuth

content exceeds 47 wt % (4). During production of bismuth bearing alloys by

conventional continuous casting the alloys would expand inside the cooled mold

and thus give rise to an increase in frictional force and hence be difficult to

produce.

With respect to replacement of lead-bearing alloys, it has been found (4) that

there are some eutectic alloys containing bismuth which could meet the required

properties.

The following features summarize the advantages of bismuth bearing alloys

for use as solders,

Bismuth bearing alloys have a wide range of rnelting temperature from 72°C

to 212°C. This is required to widen the range of applications under different

conditions.

Like lead, if bismuth is added to t h it reduces the surface tension and

improves the wetting behavior.

Sn-Bi alloys provide a better matching of thermal expansion coefficients with

a copper substrate than Sn-Pb alloy.

In the present study a ternary eutectic alloy that has the maximum proportion

of bismuth in weight % and is brittle in nature was selected from Bi-ln-Sn system.

The composition of this alloy in weight % is 57.2 % Bi, 24.8 % In and 18 % Sn

and the eutectic temperature is 77°C.

from the manufacturer point of view, the high bismuth content of this alloy

can create problems due to brittleness with conventional casting processes.

Even if this alloy could be produced, the breakage of the product during

transportation would create loss. Again if this alloy is to be used in coiled wire

form it cannot be produced by conventional methods since the increased number

of processing steps would increase the cost of production. The airn in this study

was to produce this material in one step as net shape wire with minimum

standard deviation in cross-sectional area and a smooth defect free surface. A

further objective was tu conduct a comparative study on microstructure and

mechanical properties between OCC wire and statically cast product.

Considering al1 of the factors mentioned the main objectives of the present

study are as follows,

To continuously cast a lead free ternary eutectic alloy which may be quite

impossible to produce by the conventional casting process.

To generate net shape cast wire of srnall cross-sectional area with good

surface quality, free from intemal defeds and with minimum deviation in

dimensional stability.

To make a comparative microstructural study and check segregation

behavior within OCC wire and statically cast products.

To compare the mechanical propeiiies of OCC wire and statically cast

products.

CHAPTER 2

BASIC PRINCIPLE OF OHNO CONTINUOUS CASTING

During the 1980's Professor A. Ohno of Chiba lnstitute of Technology,

Japan pioneered the development of Ohno Continuous Casting and its

application in various fields of casting (5).

In general the casting process involves the solidification of liquid metal

following different rates of heat extraction to obtain the desired shape and

properties simultaneously. The mechanical properties of the cast product depend

on the size and orientation of the grains. To obtain consistency of mechanical

properties, phases should be distributed uniformly through out the matrix. In the

early 1970's the formation mechanism of equiaxed grains was investigated

(6,7,8). It was found that equiaxed grains fomed at the mold wall at the initial

stage of solidification and were carried to the center of the ingot through

convection. These works led to the development of the Ohno Continuous Casting

(OCC) process (5).

Figure 1 shows the main difference between the OCC and conventional

continuous casting process. In conventional casting process the mold is water-

cooled and the nucleation of the crystals starts at the mold surface. The growth

direction is from mold surface toward the center of the cast strand and

perpendicular to the casting direction and thus a multi-crystalline product is

produced. Cooled mold solidification leads to formation of segregation and

shrinkage cavities. Since solidification starts at the mold surface, this is the

source of frictional force between the mold and the cast strand. For the casting of

brittle and strain sensitive material with small cross-sectional area the frictional

force betvireen mold and cast strand could be high enough to cause breakage

and interruption during casting.

A) CONVENTIONAL CASTING

COOLED

B) OHNO CONTINUOUS CASTING

- -. . CASTING DIRECTION

...

Figure 1 : Schematic diagram showing the principle of OCC and the difference with conventional continuous casting.

In the OCC process the mold is not cooled but heated externally and the

temperature of the mold is kept above the melting point of the metal to be cast.

The cooling device is positioned in front of the mold exit to extrad heat from the

cast strand in a direction parallel to casting direction. In this way unidirectional

solidification is possible and under specific casting conditions, cast products with

a long single crystal can be produœd. The distance of the cooling device from

the mold exit can be adjusted according to the casting speed. In the OCC

process the nucleation of crystals on the mold wall is avoided and frictional force

between mold and cast strand can be minimized. Reduced friction helps to obtain

net shaped products of small cross-sectional area with a smooth surface finish,

which eliminates surface defects.

Thus OCC process can improve the properties of the cast product in

several ways. A review report by Ohno (9) detailed the features that can be

achieved through OCC process as follows,

An alternative route to conventional casting processes where it may be

difficult to produce cast products from brittle materials having a small and

complicated cross-sectional area.

Production of cast products having unidirectional or single crystal

microstructure with improved mechanical properties.

Generation of net or near net shape cast products having a smooth surface

quality and freedom from defects such as shrinkage cavities and blowholes.

Elimination of the requirements for further processing and minimization of the

cost of production.

Provision of good workability.

CHAPTER 3

LITERATURE SURVEY

3.1 Lead alloys

The toxicity of lead and its alloys is well known. For example the safe

drinking water a d amendments U.S., 1986, prohibited the use of lead pipe and

lead containing solders for drinking water lines. Recently the US. environmental

protection agency regulations considered the ban of lead-containing solders (3,4,

10,ll). The increasing use of electronic equiprnent and the use of lead-bearing

solders may create a severe pollution problem not only on the industrial shop

fioor, but also during disposal and the subsequent effect on the environment.

Lead-free solders available are based on tin, indium or bismuth alloy systems.

Other elements are added to lower or increase the liquidus temperature (12).

3.2 Effect of bismuth on mechanical properties

Shewmon (13) considered the effect of bismuth on the mechanical

properties of cast gold. It was observed that an addition of only 0.2 % bismuth in

gold reduced elongation and ultimate tensile strength to zero. The fracture

surface was crystalline, which indicated brittle fracture. It was also observed that

the solid solubility of bismuth in gold was low. It was found that liquid metal

embrittlement occurred at low fracture stress with increase in bismuth

concentration.

Kariya and Otsuka (14) studied the effect of bismuth in Sn-3.5 % Ag alloy.

They found that the addition of bismuth beyond 2 % advenely affected the

fatigue life of the alloy. Fatigue life is defined as the nurnber of cycles at which

the stress is half of the maximum initial applied value. They found that addition of

bismuth increases the tensile strength due to solid solution hardening or 8

strengthening due to dispersed particles of bismuth but the ductility in ternis of

reduction in area decreases dramatically . Glazer (4,ll) emphasized the importance of physical and mechanical

properties of lead-free ailoys. Studies of solder alloys include the investigation of

melting temperature, surface tension, electrical resistivity, microstructure, and for

mechanical properties, time independent monotonic tensile strength, shear

strength and elongation as well as time dependant monotonic temperature

related creep. Homologous temperature is defined as the ratio of working

temperature to rnelting temperature in absolute scale. Since the melting point of

solder alloy is low, long exposure to a temperature which exceeds 50% of the

homologous temperature is important because it may give rise to grain growth

phenomenon which affects the creep retated properties. Thus fine microstructure

is important to obtain improved mechanical properties.

Mei and Morris (15) stated that mechanical properties are dependant on

microstructure, the combination of phases and their distribution, which depend on

several factors. Faster cooling rates yield fine microstructure. Sig nificant

microstructural differences and their effect on improved fatigue life due to cooling

rate have been observed in the case of Pb-Sn alloy. In the case of lead free

alloy, these investigaton did not mention about cooling rate, which is an

important issue with respect to the solder microstructure and its mechanical

properties.

McCormack et al (16) noted several improvements during investigation of

Bi-Sn solder alloys as a substitute for Pb-Sn alloy. Bi-Sn eutectic alloy has a low

melting temperature (139°C) and this penits the use of inexpensive circuit

board. They exarnined the importance of faster cooling rates and the addition of

grain refiner on microstructure. After quenching in ice water, they found that for

the same sample the abrupt strain rate sensitivity changed from 60 inchlinch to

10 inchlinch with increasing strain rate from O.Ol/second to 0.1Isecond. Again at

a strain rate of 0.01lsecond the total strain was 10 % and 28% for slowly and

forced air-cooled samples respectively. It is expected that the forced air-cooled

samples would have finer microstructure. It was menüoned that the strain rate

sensitivity was due to the bismuth rich phases and their continuity. Silver was

added as grain refiner, but special precautions are necessary to avoid an excess

of silver and the formation of high melting point phases.

Jin and McComack (17) examined the behavior of Bi-43%Sn eutectic

alloy and reported a "lamella like microstructure" composed of p-Sn and Bi phase

when solidified to roorn temperature. From binary diagram of Bi-Sn it can be

observed that Sn has partial solid solubility with bismuth. At approximately

100°C, in Bi-Sn system the Sn rich side shows that approximately 10% of

bismuth dissolves into the P-Sn phase. On cooling to room temperature bismuth

precipitates in coarse P-Sn phase. It also undergoes signifiant microstructural

coarsening resulting in a nonuniform distribution of phases which affects

mechanical properties. To obtain fine microstructure these authors added

dispersed particles to inhibit grain growth and serve as nucleation sites. The

condition related to the selection of such dispersoids was that there should be no

solubility or reactivity with the matrix. To obtain uniform distribution and thereby

avoid agglomeration of the dispersoids, a novel magnetic distribution technique

was proposed.

3.3 Microstructure of eutectic alloys

Depending on the morphology of the microstructure a binary eutectic

structure can be classified as either regular or irregular. In the regular class,

lamellar or rod like structures of two phases are fonned whereas in the irregular

structure one phase forms skeletal faceted crystals and the other nonfaceted

phase grows inside the first, with the formation of a complex regular structure

(18). Croker et al (19) investigateâ different morphologies of eutectic binary

alloys. They concluded that at a given growth rate if the entropy of solution and

relative volume of each phase are known the unknown structure of a particular

eutectic alloy can be predicted. Binary eutectic alloys can also be classified into

three categories depending on the nature of the solid-liquid interface (20). In a

binary system with Wo phases, three different types of interface are possible: 1 O

nonfaceted-nonfaceted, faceted-nonfaceted and faceted-faceted. The

morphology of the microstructure depends on the type of interface, local growth

condition and temperature gradient. For example, in a binary eutectic structure

with h o phases A and B, if both A and 6 are nonfaceted then the morphology

will be regular lamellar or rod like. If one of the phases is faceted and the other

nonfaceted then this will produce either irregular or complex regular

microstructure. The complex regular structure is due to the ingrowth of a

nonfaceted phase inside the faceted phase (20). This forms a skeletal structure

with different shapes which depend on the local solidification condition. In Figure

2 different morphologies of microstructures have been displayed for binary

eutectic alloys (20,21).

In the case of ternary eutectic alloys, three-phase lamellar structures have

been reported. The combination of ABCBA regular lamellar structure is found in

the Pb-Sn-Cd systern (22), Figure 3a. Pb-Sn grew in a coupled manner, but not

Sn-Cd due to the requirement of high interfacial energy due to the presence of

Cd faceted phase. Another lamellar structure was observed in one of the eutectic

compositions (melting point 332°K) of the Bi-ln-Sn system (23). It was observed

that fibers of In rich phase and Sn rich phase fonned a regular structure in a

matrix of Mnz phase, Figure 3b. However a regular lamellar morphology of

ABCABC was not reported.

Ruggiero and Rutter (24) examined the microstructure of the eutectic alloy

57.2 % Bi, 24.8 % In and 18 % Sn with a solidification temperature of 77S°C.

They used a technique of slow unidirectional growth with a growth rate of 0.74-

53mmlday and quenched the sample to reproduce the solid-liquid interface. With

slow growth rates up to a maximum of 1 .ômmtday the microstructure consisted of

two regions Biln-y Sn and massive Bi, Figure 3c. With increase in growth rate,

two regions of binary structures are found, Biln-y Sn and Bi-y Sn, as shown in

Figure 3d. According to the morphology of the microstructure the lamellar Biln-y

Sn was described as a quasi-regular structure. Biln was observed as faceted

phase and y Sn as non-faceted and according to the classification this binary

structure should be under irregular class, but the morphology is regular lamellar

11

structure of Biln and y Sn and hence termed as quasi-regular. The other Bi-y Sn

binary structure was a complex regular structure due to the faceted bismuth

phase. In both binary structures y Sn, common to both binary phases, is non-

faceted. It was suggested that the hexagonal ySn phase decornposed below the

eutectic temperature by a ternary eutectoid reaction to form PSn, Biln and Bi

phases. In the case of Bi-y Sn complex regular structure, bismuth of decomposed

structure was not present. It was suggested that in the case of Biy Sn complex

regular structure, bismuth has either diffused or been incorporated inside the

bismuth phase.

In the Bi-Cd-Sn system Ruggiero and Rutter found that the three phases

formed were Bi, Cd and P Cd-Sn. P Cd-Sn is a hexagonal Sn rich phase of Cd-

Sn binary system. The growth rates, frorn 8 nmlsec to 1.1 pm/sec, were used

and the three phases were not an intimate mixture (25). At a higher growth rate

the morphoiogy was, a quasi-regular or complex regular binary structure of Bi

and p Cd-Sn, outlined by Cd flakes. At slower growth rate a regular lamellar

microstructure of Cd and PCd-Sn was often formed with large and irregular

masses of Bi. It should be noted that pCd-Sn is non-faceted and common to both

binary structures. The evidence of ternary eutectic decomposition was also

observed in the Bi-Cd-Sn system where pCd-Sn decomposed into Bi, Cd and a-

Sn.

Also in the Bi-Cd-Pb system Ruggiero and Rutter have referred to regions

of double binary structure with the formation of a quasi-regular binary structure of

Bi and Pbz6i (26). The second binary structure consists of broken larnellar Cd

and Pb2Bi, where PbzBi is a non-faceted phase and common to both binary

structures. The growth rate range in this study was from 10.6 nmlsec to

555nmlsec.

In general it was observed that in the case of double binary structures one

nonfaceted phase is common to both binary structures. All the eutectic alloys

reported were solidified under very siow growth rates in the range of 5.7 nmlsec

to 53 mmlday and there was no report related to higher growth rate.

In some bismuth bearing alloys the formation of massive blocks of bismuth

has been reported when the slowest growth rate was maintained depending on

the binary or temary system used (20,23,24-28). For example (24) in the slowest

growth rate of 0.74mmlday and 1.6 mmlday, the bismuth was found in block

shape. When the growth rate was beyond 1.6 mmlday complex regular structure

of bismuth with y Sn was fomed. It was suggested that the complex regular

structure was due to the presence of a faceted bismuth phase. All directional

solidification experiments were carried out under controlled temperature

gradients and at very slow growth rates. It is expected that due to the higher

density of the Bi phase than bulk liquid that bismuth would segregate, but there

was no report of this.

The growth of dendrites in eutectic composition has been observed

(27,28) when the alloy system has a skewed coupled zone below the eutectic

temperature.

3.4 Segregation

Segregation results in the non-unifonn distribution of different phases and

affects consistent mechanical behavior. Segregation is the result of 1 ) rejection of

solute at the solid-liquid interface during solidification and its distribution by

diffusion and mass flow away from solidification front, 2) nucleation of primary

phase and segregation depending on the density of the primary phase with

respect to the bulk liquid. For example segregation was reported for Sn-Pb

eutectic alloy (29) in which primary lead dendrites nucleated in the undercooled

rnelt and segregated to the bottom of the ingot whereas tin rich dendrites

segregated to the top.

Grugel (30) studied a macrosegregation phenornenon in lead-tin alloy

during unidirectional solidification. Two alloy compositions, Pb-45Sn and Pb-

75Sn, were selected. In the case of Pb-45Sn hypoeutectic alloy the primary lead

dendrites rejected less dense tin rich eutectic. In vertical casting the tin rich

13

eutectic was found between the Pb dendrites and in the case of horizontal

casting the lin rich eutectic phase segregated to the top since it was less dense.

For the other alloy, Sn was the primary dendrite phase and the denser lead rich

phase segregated to the bottom in horizontal casting. Microstructural

homogeneity was absent in both cases due to segregation. To obtain uniform

distribution of the segregated phase, different positions and furnace rotation were

checked. The most suitable condition to minimize the segregation was obtained

by an axial rotation of the furnace at 10 rpm and positioning it at an angle of 5"

from the horizontal position. However there was a lack of technical information

and this appeared to be a trial and error method. No suggestion was made to

minimize segregation where gravitational force is involved.

Segregation in continuous casting of steel and the adverse effect on

mechanical properties, Le., tensile, fatigue and impact toughness has been

reported (31,32). The mass movement is greatly influenced by convectional force

and depends on solute concentration difference, temperature difference in the

liquid and gravitational forces on crystals growing in the liquid. It has been

observed that the segregation of carbon, sulphur and to a lesser extent

manganese and phosphorus has resulted in non-uniform mechanical properties.

It was found that segregation tened as centerline segregation increases with the

increase of the columnar region. This segregation which decreases the reliability

of continuously cast steel is minimized by electro-magnetic stirring methods

(EMS). This promotes the formation of equiaxed grains by increasing the

convective flow which promotes the separation of dendritic tips. The detached

dendritic üps settle within the liquid pool and act as nuclei for the formation of

equiaxed crystallization . The purpose is to d istribute rejected solute unifomly

throughout the entire structure and thus obtain more uniform mechanical

properties.

Lamellar CU-AI eutedic (21)

a) Reguîar microstructure

As polished (nat etched) 200 X

3 105 Scctlon of a castine show- the N-Si eufecttc, which consists of

short prrrtlcles of siilcon tdark) la rn aluminum matrb. Some particles ue connected in the plaac shown; othus are con- nectai in other plones. (Red 1)

As polished (not etched) 400 X Same casting as in 3097, but a sec- 3098 ,* ,,, ,&, , ,muon

of solldîficatlon. showing the c h u l u cmss section of the dark, fibrous putlclcs of MnSb phase In the antimony xnatrix.

Rod like Mn-Sb eutedic (21)

Needle shaped AI-Si eutedic (21) Cornplex regular BiSn euteaic (20) b) Imegubr microstructure

Figure 2 : Different morphologies of binary eutecüc alloys.

a)Lamellar structure in

temary eutedic (22)

Pb-Sn-Cd b) Lamelhr structure in Bi-Sn-ln

temary eutectic (23)

c) Quasi-reg ular lamellar strudure

with blocks of bismuth in Bi-Sn-ln

temary eutectic at slow growth rate

(24).

d) Double binary structure composed of

quasi-regular and cornplex regular

regions in Bi-Sn-ln temary eutecüc at

faster growth rate (24).

Figure 3 : Different morphologies of temary eutectPc alloys.

CHAPTER 4

EXPERIMENTAL ASPECTS

4.1 Alloy preparation

The composition of the alloy was selected from one of the ternary

eutectics within the Bi-ln-Sn system as shown in Figure 4 (24). The composition

in weight percentage is 57.2% Bi, 24.8% In and 18% Sn and the melting point of

the alloy is 77S°C (33). All the individual metals were 99.99% in purity as

received from the supplier. Bismuth was received in tear shaped grains with a

dull surface appearance due to the presence of oxide. To get rid of oxide the

bismuth grains were melted in a graphite crucible in the presence of argon gas.

The oxide was removed from the top of the melt and the clean melt was poured

into a shallow graphite mold. Tin was received as a large block and melted in the

same way to obtain thin strip. Indium was received in a sealed condition.

Bi, In and Sn were weighed according to the exact proportion for the alloy

composition and placed in a graphite cylindrical crucible. The crucible was

heated extemally by propane torch and argon gas was supplied from above to

avoid any oxidation during melting. The mixture was stirred with a glass rod to

ensure homogeneity. To obtain small pellets to facilitate feeding during casting

and avoid segregation, the melt at a temperature of approximately 95' C was

poured in a stainless-steel vessel containing water. The solidified pellets were

rernoved from water and dried at room temperature. The composition of the alloy

was confirmed by neutron activation analysis and melting point measurements.

Figure 4: Liquidus projection of Bi-ln-Sn system shdng 350.5 K

temary eutectic (24).

4.2 Equipment

Figure 5 shows a schematic diagram for producing wire by the OCC process.

The melting and casting apparatus was constructed from a 37mm long and 30

mm diameter graphite rod. An open cavity at the middle of the graphite rod was

made to hold molten metal and to melt the alloy. The dimension of the cavity was

16mm wide, 25mm long and 20mm deep. A channel was drilled at one end of the

mold cavity length 4mm and diameter 2mm.

A heating element was prepared from nichrome wire of 0.45mm diameter.

The wire was coiled around a graphite crucible coated with alumina cernent for

electrical insulation and then protected by another layer of alumina to avoid

mechanical damage during handling.

The control K-type thenocouple was calibrated with boiling water and an ice-

water mixture and a correction factor was determined. A hole 4mm deep was

drilled on top of the graphite crucible near the mold cavity channel at a position

1 mm away from the channel entrance.

In front of the mold exit a cooling device was attached to a movable platform

to adjust the position from the mold exit. The cooling device consists of a pipe

having 2mm inner diameter and connected to a water supply source. A stopper

valve fixed in between the delivery and supply lines controlled the flow rate of

water.

A stainless-steel pipe 2mm outer diameter and 1 meter long was used as a

dummy bar to initiate freezing from the channel at the beginning of the casting. A

steel wire ?cm in length and 1mm in diameter was attached to one end of the

dummy bar to insert it smoothly inside the channel. The pinch rollen were

connected with a gear and chain mechanism and this arrangement was

connected to a variable speed control motor, so that withdrawal speed of the cast

product could be controlled.

HEATER THERMOCOUPLE

WATER COOLER LlQUlD METAL

- CASTING DIRECTION

INSULATOR I

SOLID- LlQUlD INTERFACE

Figure 5 : Schematic diagram of OCC equipment for generation of net shape wires.

A horizontally placed platfon was positioned in between the water cooler

and pinch roller and the dummy bar was passed through a glass tube fixed on

top of the platform. In this way mechanical vibration dumg casting was reduced.

4.3 Experimental procedure

The graphite crucible and heater, was aligned to ensure that the mold cavity

channel was at the same level as the dummy bar and pinch roller.

The dummy bar was inserted through the pinch roller and then through the

glass tube on the movable platform. The steel wire connected with the dummy

bar was introduced inside the mold channel. The dummy bar was located

between the rollers with the help of a manual turning screw on the top rollers. An

adjustable knob attached to the control panel was used to regulate the speed of

the motor.

The water Row rate was fixed at 150mL per minute. The rate was maintained

the same for different casting speeds. The furnace was switched on and the

control thermocoupfe temperature adjusted to obtain a melt temperature of 87°C

at position B (Figure 5), during casting. Small pellets of alloy were fed inside the

moid cavity. In the liquid state the height of the metal head was 6mm from the top

of the channel entrance at the start of casting. During casting the metal head

height was maintained within the range of 3-6 mm to ensure appropriate metal

pressure at the mold exit.

The speed controller was set to the desired casting speed and the distance

behnreen the mold outlet and the cooler adjusted to bring the solidification front

just outside the mold exit.

LENGTH (mm)

. -- - - - .

I

* t * l

Length inside mold cavity Solid-liq uid interface

Length of mold channel

Figure 6 : Temperature profile during casting.

When the melt temperature was 87°C at the channel entrance, a temperature

profile was measured as shown in Figure 6. The temperature was measured

from point A to B inside the mold cavity at a distance interval of 5mm. Then the

temperature was measured at the mold exit and the temperature at the solid-

liquid interface was considered to be the solidification temperature of the alloy.

The control themocouple temperature was recorded as reference to follow the

particular temperature distribution.

During casting of wire 2mm in diameter, with the above temperature profile,

the most suitable distance between the mold exit and cooler was 4mm and 5mm

for casting speeds of 79mm/min and 14mm/min respectively.

4.4 Static casting facility

Statically cast rod was produced inside a glass tube, 3.2mm inner

diameter. A hollow cylindrical furnace blocked at one end was prepared as

shown in Figure 7. A nichrome wire 0.45mm diameter was coiled around the

alumina cement protection over the furnace. Then another layer of alumina

cernent was applied to cover the coil. The purpose was to insulate the heating

coi1 and protect it from mechanical damage. A control themiocouple was

positioned in a hole drilled about Imm away from point B and 4mm in depth. The

glass tube was positioned at the center of the fumace and kept horizontal. The

control thermocouple temperature was adjusted to obtain a minimum

temperature of 87°C at point B of the glass tube to ensure the alloy was in the

liquid state. When the temperature was stable, a temperature profile was

measured at intervals of 5mm from point A to point 6 as shown in Figure 8. From

the plot of temperature vs. distance it was found that from point A the

temperature was essentially constant for a distance of approximately 60mm. This

was designated as the plateau region. From 60mm to point 6 a temperature

gradient reg ion exists. Unidirectional solidification was expected in the

temperature gradient reg ion.

To perform mechanical testing the minimum length of sample required

was 100mm and efforts were made to maintain the same morphology of structure

along the entire length of the sample. A problem associated with the glass mold

was that there was no gap between the cast metal and the glass wall due to

negligible thermal contraction of the alloy since the alloy has a high bismuth

content and bismuth expands when solidified. To remove the specimen, the

glass mold had to be broken and the specimen also broke into pieces. Thus the

length of the specimen was insufficient for tensile tests. Difterent structural

rnorphology in the plateau and temperature gradient regions was expected and

the results will be discussed later. To resolve the breakout problem a graphite

split-mold, Figure 9, was designed.

HEATER ALUMINA CEMENT f HERMOCOUPLE

BLOCKED END

Figure 7 : Schematic diagram of resistance fumace for static casting in a glass mold.

I PLATEAU REGION 1 GRADIENT REGlOl

90 - GLASS MOLD

80 -

LENGTH (mm)

Figure 8 : Temperature profile inside the glass tube showing a plateau region and a temperature gradient region.

The split mold helps to remove the specirnen in a smooth manner without

breaking the specimen and the desired length of specimen can be maintained.

60th ends of the split mold were connected by threaded graphite caps to hold the

split mold firmly in place and to avoid liquid metal leakage. The diameter of the

mold cavity and length are 3mm and 1OOmm respectively. The extemal diameter

of the split rnokl was close to the inner diameter of the resistance furnace in

order to reduce the air gap and enhance better heat flow between the furnace

and graphite split mold, Figure 10.

A small hole was drilled parallel and equal to the length of the cavity to

insert and position a thennocouple. The distance between the hole and cavity

was Imm. The temperature was measured within the hole and within the cavity

and it was found that the temperature difference for equivalent positions was

negligible. In this way the temperature of the melt can be detemined from

measurements within the hole without inserting the thennocouple inside the melt.

The control therrnocouple temperature was adjusted to obtain a minimum

temperature of 87'C near point B. When the temperature reached 87°C a

temperature profile was measured from point A to point B at intervals of 5mm as

shown in Figure 11. It can be seen from the profile that the temperature is

essentially constant (Maximum difference i 1°C) along the entire length of the

mold.

Figure 9 : Photograph of graphite split mold.

I BLOCKED END OF MOLD

THERMOCOUPLE HOLE

SPLIT MOLD

MOLD CAVITY

CONTROC TC

.. .-. . ... . . ...

.. ......'... *;*...,,* ...

THREADED CAP OPEN END

Figure 10 : Schematic diagram of resistance fumace for static casting in graphite mold.

... --.

MOLD CAVITY HOLE l

SPLIT MOLD

LENGTH (mm)

Figure 11 : Temperature profile inside the mold cavity.

4.5 Experimental procedure for static casting

The fumace was first positioned vertically keeping the blocked end at the

bottom to support the mold and to facilitate feeding of OCC wire material through

the top open end of the glass rnold. A themocouple was positioned at the middle

of the glass tube to measure the cooling rate and to confinn the temperature of

the alloy. The furnace was switched on and when the control thermocouple

temperature was stable to give a minimum temperature at point 6 of 87'C the

OCC wire material was added. When the solid metal was converted to liquid

state the furnace was positioned horizontally. The horizontal position was

selected since OCC casting is also horizontal. The furnace was switched off after

3-4 minutes and the temperature drop was recorded at regular time intervals.

When the temperature dropped to approximately 50°C the glass tube with

solidified alloy was removed and cooled to room temperature.

The specimen was removed by breaking the glass tube and in most cases

the specimens broke into pieces. The specimen pieces were marked properly to

differentiate them from plateau and directional region and also the bottom and

top part to prepare samples for microstructure.

In other experiments samples were produced for tensile tests using the split

mold. The purpose was to obtain samples 1OOmm in length with the same

microstructural morphology along the entire length and cross section of the cast

products. Special precaution was taken to avoid any flow of liquid metal through

the gap of the split mold during casting. The themocouple was positioned about

50mm away from point A to obtain a cooling curve and confinn complete

solidification of the alloy. At the beginning, the furnace was positioned vertically

and OCC wire was fed from the top. When the metal inside the mold cavity was

in the liquid state the top cap was fixed and the thermocouple was inserted. The

furnace with the split mold was then positioned in the horizontal location. The

furnace was switched off aRer 3-4 minutes and when the temperature reached

50°C the split mold was taken out and cooled to room temperature. The threaded

caps from both ends of the mold were opened and the cast sample removed. Ali

sarnples produced were stored in a cardboard socket to prevent buckling and

bending that could cause centering problems during machining.

4.6 Sample grip design for mechanical testing

INSTRON machine, mode1 8501, with ASTM specification k 1 % was used

for the tensile test to compare the ductiiity of the OCC and statically cast

specimens. A photograph of the machine is shown in Figure 12. The machine is

equipped with two grips and can perform tensile or compressive tests under

different controlled conditions.

Problems were encountered with the l NSTRON machine as follows,

a) The grip was flat and to perfon tensile tests with round samples, the area of

contact was from two points as shown in Figure 13 A. Due to insufficient area of

contact, the specimen would slip and the data on elongation % was wrong.

b) The grip was designed to operate with a high hydraulic pressure. The high

impact and the compressive force deforrned the specimen at the grip area. It was

found that OCC samples became thin strip at the grip area as described in

Figure 138 and static samples originated cracks around the surface and were

deformed in the grip area as shown in Figure 13C. The result was premature

failure in the deformed area and not inside the gauge length and again the data

from testing was wrong.

c) For a tensile test it is important to align samples along the stress axis

otherwise the test could give wrong information. With the INSTRON machine the

problem was that there was no proper marking on the grip to align the sample

along the stress axis.

Figure 12 : Photograph of INSTRON machine, mode1 8501.

Two point pressure on ( ) a round specirnen

A) Specirnen in INSTRON flat grip without pressure

Ductile specimen

1

6) OCC -*men at grip area plastically defmed after pressure

C) Static specimen at grip area deforrned and cracked

Figure 13 : Schematic diagram showing the problem associateci with INSTRON high pressure and flat gnp

To eliminate these problems a special grip was designed as shown in

Figure 14. Two pairs of low carbon steel plate with dimensions 3mm thick, 41mm

width and 56mm length were machined. The dimensions were important to

position each pair of modified grips between the lower and upper grip of the

INSTRON machine to avoid play inside the INSTRON grip. A "V" shaped right-

angled groove was cut at the center of each plate along the length to position

samples perfectly along the stress axis. The "V" shape was selected to obtain the

cross-sectional shape of a square after assembling two plates. The distance

between the two sides of the square was calculated to be exactly equal to the

diameter of the specimen. In this way the holder groove touches the specimen

from four points as shown in Figure 15 a. This helps to maximize the contact

area and slipping during the tensile test can be avoided. Again the hydraulic

pressure of the INSTRON grip was resisted by the rnodified low carbon steel

grips.

To carvout experiments with samples of different diameter, it was

necessary to prepare holders with different dimensions of the right-angled "V

shaped groove. If the diameter of the specimen is "2X, then the depth and width

of the groove can be calculated from Figure 15 b and is shown in tabulated form

in Table 1.

Table 1 : Mathematical expression to obtain the dimension of the groove of the

modified grip according to diameter of the sample

Diorneter of sample :

Depth of the groove Width of the groove

INSTRON flat gflp

Modified grip Specirnen

a) A transverse sedion of modified gnp showing the four point contact and resistance to high pressure from INSTRON machine

Diameter of specimen - 2X, Height of the groove - X 4 2. Wdai of the gmove - 2X 4 2

b) Dimension of rnodified grip groove according to the diarneter of the specimen

Figure 15 : Schematic diagram of modified grip.

4.7 Sample preparation for tensile test

The wire produced by OCC has good surface finish and during machining

there was no problern associated with centering of the specimen and to obtain a

constant gauge diameter. In contrast the surface of the staticalîy cast product

was not smooth. The surface quality can be distinguished from the photograph in

Figure 16. The surface of the statically cast product has blowholes from

entrapped gas and ridges which originated from the gap between the splits.

These specimens were machined until the surface was smooth enough along the

gauge length to eliminate surface defects. The depth of shrinkage cavities varied

from sample to sample.

The static cast product was prone to breaking if the depth of cut exceeded

a certain critical value per pass. The depth of cut per pass was fixed within the

range 0.i mm to 0.1 5mrn per pass and this helped to avoid not only breaking but

also buckling and deflection of the samples. For OCC samples there was no

problem with surface defects or breaking of samples. To maintain the same

surface condition after machining, the same depth of cut per pass was

maintained. For both static and OCC samples special precaution was taken to

avoid any rise in temperature for the low melting point alloy by ensuring sufficient

fiow of cutting fluid.

a) OCC wire without surface defect

b) Statically cast rod with surface defect

Figure 16 : Photograph of OCC wire and statically cast products showing

significant difference in surface quality.

4.8 Sample preparation for microstructural observation

Both OCC and statically cast samples were cut into small pieces and

placed in a plastic cup. Epoxy resin 8 parts and hardener 1 part were mixed for at

least 2 minutes and poured into the plastic cup on top of the sample. The amount

of hardener was limited to 1 part by volume to avoid an increase in temperature.

The mounted samples were polished with emery papers from coane to fine

following the numbers 320, 500, 800, 1200 and 2400 respectively. During

polishing, sufficient water fiow was maintained to avoid rise in temperature due to

friction, considering the low melting point of the alloy. Final polishing was carried

out with alumina paste of lpm size and at the end with 0.3 pm size.

For image analysis samples were etched. For compositional analysis,

etching was not canied out to avoid dissolution of elements from the samples.

For both purpose the sample was either gold or carbon coated to obtain better

resolution. The etchant used was "Rhine's Etchantt'(34) which has the following

composition:

Hz0 - 300ml. K2Cr07 - 6gm, H2s0.1- 201111, NaCl - 12ml (saturated solution),

1 part of etchant and 9 parts of water by volume were mixed to dilute the

etchant. A cotton swab was dipped into the diluted etchant and the sample

surface was covered. The etching time was 30 seconds and after etching, the

samples were washed first with water and then with ethyl alcohol. At the end the

samples were dried with compressed air.

CHAPTER 5

RESULTS AND DISCUSSION

5.1 Evaluation of casting condition for OCC

5.1.1 The occurrence of breakout

Experiments were conducted at different casting speeds to determine the

most suitable casting condition for Bi-ln-Sn ternary eutectic alloy (Melting point

77°C). The purpose was to avoid run out of liquid metal outside the mold exit

during conünuous casting.

In the OCC process the mold temperature is kept above the melting point

of the metal to be cast and therefore solidification does not occur inside the mold.

The liquid metal is cooled by water in front of the mold exit and heat is extracted

through the cast strand. Thus solidification occurs between the mold exit and

water cooler. The liquid rnetal outside the mold exit is a floating zone held in

place by surface tension. Breakout occurs when the volume of the liquid rnetal

overcomes the surface tension. In addition to the surface tension, the volume of

the floating zone is influenced by several factors like water-cooler distance, flow

rate, casting speed, mold exit temperature, height of the metal head and cross-

sectional area of the products.

In the experimental procedure some controlling factors which influence

breakout were kept constant. The melt temperature at the channel entrance was

kept constant at 87'C and the height of the metal head was kept 6mm maximum

from the top of the mold channel for al1 casting speeds. The purpose was to

check the effed of water-cooler distance at different casting speeds. At a casting

speed of 14mmlmin the critical water cooler distance was 5mm away from the

mold exit. It was observed that if the distance was more than 5mm the casting

operation became sensitive and any small vibration caused bulging and

subsequent run out of the liquid metal. At a casting speed of 79mmlmin the

critical water cooler distance was 4mm.

Another parameter to control the breakout of metal is the height of the

metal head. An increase in the height of the metal head will increase the

pressure at the mold exit. Consequently the flow of liquid metal will increase and

it becomes difficult to keep in hanony with the rate of solidification and the

supply of liquid metal. The result is bulging and ovefflow of liquid metal at the

mold exit and if the heat extraction is not enough the liquid metal will run out.

During casting it was obsenred that if the metal head was more than 6mm from

the top of the channel the diameter of the wire increased above 2mrn and this

could interrupt the casting process. If the metal head height was less than 3mm

at a casting speed above 79 mmlmin the diameter of the wire was reduced to

less than 2mm due to insuffcient metal pressure and Row at the mold exit. So it

is important to keep the height of the metal head constant to keep the metal

pressure at an optimum value at different casting speeds to avoid any

discontinuity of the casting process and fluctuation in the cross sectional area of

the product.

5.1.2 Surface appearance of cast wire

In the present work the bismuth content of the eutectic alloy was 57.2%.

There are problems associated with high bismuth content, one is the alloy

becomes brittle in nature and the other is the alloy becornes a poor conductor of

heat. McCorrnack et al (16) discussed the severe strain rate dependant

embrittlement of the Bi-Sn solder due to the presence of bismuth rich phases.

Special precaution was taken during the present work to avoid mold strand

friction by avoiding solidification inside the mold. The solid-liquid interface

position can be controlled outside the mokl exit by positioning the water cooler

distance and controlling the mold exit temperature (35.36 and 37). If the water

cooler distance is too short the heat extraction will be high enough to shift the

interface inside the mold. If the water cooler distance is kept constant and the

mold exit temperature is below the critical temperature for a particular constant

casting speed, solidification will occur inside the mold. Under some conditions if

solidification starts well inside the mold, the frictional force between mold wall

and the cast strand will cause rough surface and if the frictional force is

excessive the casting operation will be interrupted.

In the present study the moldçhannel entrance temperature was

maintained at 87"C, water fiow rate at 15Omllmin and feed rate was controlled to

keep the metal head below 6mm. The different surface appearance was

observed by varying the mold-cooler distance. At a casting speed of 14mm/min

when the mold-cooler distance was below the critical distance of 5mm, it was

observed that the surface of the wire was not smooth and breakout occurred due

to friction between the mold wall and cast strand. Again at a casting speed of

79mm/min, when the mold-cooler distance was below the critical distance of

4mm, the same phenornenon was obseived. An attempt was made to cast at

145mmfmin but the critical mold-cooler distance was then below 2mm and the

inconsistent overllow of water was nearly touching the mold exit. The result was

43

a temperature drop at the mold exit and solidification inside the mold. The critical

casting parameters are indicated in Table 2 for different casting speeds.

In initial experiments, the appearance of ridges at regular intervals on the

cast surface was observed. The reason for this was mechanical vibration from

the pinch rollers that caused disturbance at the solid-liquid interface. Other

problems associated with mechanical vibration were a wavy appearance of the

wire and a lack of straightness. lt was found that the pinch rolls were not aligned

properly and this created an up and down movement of the wire and

consequently on the liquid metal at the mokl exit. The result was production of a

wavy surface. To avoid this the pinch rollers were aligned properly, the wire in

between was pressed firmly under load and passed through a glass tube fixed

with a clamp against a horizontal platform.

With the casting conditions shown in Table 2, the OCC wire produced had

a smooth surface finish with no trace of casting defects.

Table 2 : Optimum casting conditions for Bi-ln-Sn alloy

/ Cooling wsbr flow rate (mllrnin]

Temperature at mold

channel enûance ( O C ) -

Water-cooler distance

1 From mold exit (mm)

Casting speed

14mmlmin.

Casting speed

79mm/min.

Casting speed

14Smmlmin.

5.2 Microstructure

5.2.1 Microstructure of static cast rod

From the temperature profile of the glass tube in Figure 8 it can be seen

that from position A to 60mm the region is a plateau and from 60mm to position B

there is a temperature drop and the region has a temperature gradient. So it was

expected that the microstructure would be different in the two regions. The

samples were selected, one from the center of the plateau region and the other

from the center of the temperature gradient zone. The temperature drop or

cooling curve of the melt in the glass tube was measured, one at the middle of

the plateau region and another at the middle of the temperature gradient reg ion.

The cooling rate was 0.7'CImin for both regions as shown in Figure 17. Frorn the

cooling curve, evidence of a clear plateau region, which is the freezing point

(77.5"C) of the ternary eutectic alloy, can be seen. No inflection temperature

other than 773°C was observed in the cooling curve. This indicates typical

cooling be havior for a eutectic system.

Figure 18 and 19 show the microstructure of static cast rod frorn the

plateau reg ion and tempe rature gradient reg ion respectively . The microstructure

indicates a polycrystalline structure although the composition of the alloy

corresponds to one of the ternary eutectic points in the Bi-ln-Sn system and the

plateau region of the cooling curve is typical of a temary eutectic. There is no

evidence of a regular repetitive arrangement with the intimate mixture of the

constituent phases. It can be seen from the microstructures from both plateau

and gradient region and phase identification through energy dispersive X-ray

analysis (EDX) that the microstructure was composed of several regions of

different phases as shown in Figure 20.

a) Longitudinal section

b) Transverse section

Figure 18 : Backscattered SEM image of sample from the plateau region of cast

product in glass tube showing the phenomenon of segregated blocks

and complex regular structure of bismuth (white phase).



Figure 19 : Backscattered SEM image of sample from the temperature

gradient region in glass mold showing directional solidification

and less segregation.

(cl (a Figure 20 : Energy dispersive X-ray (EDX) spectra showîng,

a) Segregated bismuth ,

b) Gray phase composed of bismuth and indium inside mottled region of

bismuth corn plex reg ular structure, c) Black dendrite composed of th with a trace of bismuth,

d) Gray spine phase inside eutecüc cell composed of bismuth, indium and

tin,

The phases identified in the plateau region were eutectic grains, black Sn

rich dendrites, white bismuth blocks and branched white bismuth structure

(corn plex reg ular structure). The growth of dendrites in alloys having eutect ic

compositions may occur due to the presence of a skewed couple zone

(27,30,38). The formation of Bi bearing phases caused extensive segregation.

In many systems the microstructure is not the intimate mixture of the

constituent phases as reported by Ruggiero and Rutter (25). In the temperature

gradient region the microstructure of the longitudinal section, as shown in Figure

19 a, is unidirectional. From both longitudinal and cross-sectional views in Figure

19 a and b respectively. there is very little segregation of the white bismuth

phase. However in the plateau region as shown in Figure 18 a and b, for

longitudinal and cross-sectional area, blocks of white pure bismuth and massive

cornplex regular structure of bismuth have segregated to the bottom part,

compared to the upper part, of the cast rod. Apart from the pure bismuth block,

bismuth is present in different shapes of complex regular structure. The white

bismuth, complex regular structure, was also found around grain boundaries. The

blocks of bismuth was identified by SEM energy dispersive X-ray (EDX) analysis

and it can be seen from Figure 20 a, that the phase is pure bismuth and there is

no trace of other elements. The existence of bismuth blocks as well as complex

regular structure segregated at the bottorn is due to their higher density. It was

found that the density of the ternary eutectic is approximately 8.5 gm/cm3

compared to that of bismuth 9.8 gmlcm3. In contrast, in the temperature gradient

region the complex structure of bismuth is along the directionally solidified

eutedic grains and the amount of segregation at the bottom is negligible. The

phenomenon of segregation has been reported for a Sn-Pb alloy (30) eutecüc

composition where Pb primary dendrites segregated at the bottom and the Sn

rich dendrites segregated at the top due to the phenomenon of difference in

density of the pn'mary phase compared to the bulk liquid. Pb had higher density

than the bulk liquid and hence segregated at the bottom and Sn had lower

density than the bulk liquid and floated to the top. This type of phenomenon gives

rise to compositional variation from top to bottom of cast products and in the final

microstructure a mixture of different phases can be found rather than the eutectic

composition. Large complex structures of Bi were observed mostly in the

temperature plateau region where the growth is not directional.

An enlarged portion of the eutectic grain of the plateau region in Figure 21

shows a tin rich black lamellar area with a trace of bismuth white phase, a spine

shaped gray matrix, mottled region of Sn rich dendrite and white bismuth

corn plex regular structure segregated around the grain boundary. The gray

matrix of the eutectic cell was identified through EDX analysis as shown in

Figure 20d, and al1 three elements, Bi, In and Sn were present. Even in the

temperature gradient region the small cornplex regular structures have

segregated along grain boundaries. The complex regular structure is the result of

the ingrowth of the second phase inside bismuth phase and the extent of it

depends on growth rate (20). In the microstructural investigation for bismuth

bearing ternary alloys it has been reported that at the slow growth rate, massive

bismuth is forrned and at faster speed, Bi phase tends to branch out to fom a

skeletal or complex regular structure (23,25). The same phenomenon was

observed by Bagheri and Rutter (20) in Bi-Pb and Bi-Sn binary eutectic systems.

The presence of three types of complex regular structures were observed

in statically cast rod within the plateau region. The three different types of

corn plex reg ular structures can be disting uished from the basic shape of "cubic",

'Yish spinen and "trigonal" as shown in Figures 23 a, b and c respectively.

Baragar et al (39) observed that the different complex regular structures depends

on growth rate and temperature gradient. They plotted growth rate vs.

temperature gradient to show at which different shapes of complex regular

structure in bismuth and lead eutectic alloy occur and found that the shape of the

complex regular structure will change from "fish spinen to "trigonal" to "cubic" with

increasing growth rate. Ruggiero and Rutter (22) found that there is a well-

defined faceted phase in Bi rich areas and to a lesser extent in the Bi-ln phase.

Hunt and Jackson (18) predicted on the basis of smooth atomic attachent at

the solid-liquid interface that if one of the phases grows as faceted, the formation

of complex regular structures should exist. Hunt and Hurle (40) stated that the

lateral movement of complex regular growth occurs along a crystallographic

plane where there should be no formation of new solid layer on an existing solid

in the presence of an energy barrier. Also Bagheri and Rutter (20) suggested that

the branching of Bi phase is due to the ingrowth of the second phase inside the

Bi faceted phase.

From Figure 23 it can be observed that inside skeletal Bi structure there

exist mottled regions. In Figure 23 c it is clear that inside skeletal bismuth the

mottled region is composed of a black and gray phase. The same mottled region

can be seen in Figure 21 in the dark dendritic region but white Bi particles are

also present. The gray phase of the rnottled region in both cases was checked

through EDX analysis as shown in Figure 20 b and it can be observed that there

is no trace of Sn phase. The black phase in the mottled region is composed of Sn

with a trace of Bi as shown in the EDX analysis in Figure 20c. The overall

composition of the rnottled region in both Figure 21 and 23 c as identifieci in the

EDX analysis in Figure 22 is identical to that reported by Ruggiero and Rutter

(24). From the phase diagram in Figure 4 it can be seen that at the ternary

eutectic temperature the tin phase should be y-Sn. From the prediction of

Ruggiero and Rutter (24) it can be suggested that pSn eventually decomposes

into p-Sn, Biln and Bi phases due to the ternary eutectoid reaction and for this

reason the mottled appearanœ of the Sn dendrites can be observed. The black

phase in the mottled region is the p-Sn phase with a trace of bismuth and the

grey phase is Biln without a trace of tin. f he solid state eutectoid reaction within

the y-Sn dendrites has occurred at different nucleation sites resulting in the

formation of irregular lamellar structure, temed "mottled" region. It can be

observed that when the structure is massive it exhibits a mottled structure and

consists of three phases (Sn, Biln and Bi). But as can be seen from an enlarged

microstructure of eutectic grain Figure 24, in some areas where the black phase

is narrow, the gray phase is absent. From Figures 21 and 24 it can be seen that

in some portions the mottled region is absent where the interphase region has

less spacing, or the microstructure is fine in that particular phase. To investigate

52

the eutectoid reaction and the branching of Bi, samples were produced by static

casting in a graphite mold at different cooling rates keeping al1 other conditions

the same. The cooling rate was maintained at 09"Clmin, O.S°Clmin, I0C/min and

2"CImin. Figure 25 and Figure 26 show the microstructures observed at

different cooling rates. It can be seen that there is no signifiant difference in

microstructure in the cooling rate range of O.S°C/min to 2"Clmin. At al1 cooling

rates it can be seen that Bi is in cornplex regular fom except in the

microstructure in Figure 25 a where it can be observed that Bi is almost block

shape when the cooling rate was Oî°Clmin. It can be concluded that at lower

cooling rates the branching out tendency of Bi can be minimized which has a

direct effect on minimizing growth rate. Advantage was taken of coarse grains at

the slowest cooling rate to investigate the occurrence of eutectoid reaction in the

primary Sn-dendrites. It can be seen from Figure 27 that the Sn-dendrites

consist of several colonies or cells having lamellar microstructure. The lamellar

microstructure in the dendrites is a clear indication of eutectoid reaction and that

each colony has grown from an individual nucleation event. It can be concluded

that when the interphase spacing is larger, the eutectoid reaction can be

confirmed by the mottled appearance with the presence of three phases Sn, Bi

and Biln. When the interphase spacing is narrow the mottled appearance

disappears and a binary structure of Sn and Bi with a regular spacing with Biln

occurs. Hence it can be suggested that when the interphase spacing is narrow

the Biln phase is incorporated into the matrix (Biln). This type of phenomenon

can be predicted when the cooling rate is faster and the microstructure is fine,

creating a narrow interspacing.

According to the findings of Ruggiero and Rutter (24) and from the

observations from the present experiments, 1 can be concluded that there exists

regions of double binary microstructure rather than temary eutectic structure

which represents the repetitive arrangement of three phases (41). In the case of

Bi-Cd-Sn system it has been observed that the Cd and P Cd-Sn phases fonned a

nearly regular binary lamellar microstructure containing coane Bi structure (25).

In general it can be seen that in the Bi-ln-Sn system double binary structures

exist in the form of Biln-ySn and Bi-ySn. ySn has the decomposed structure which

has tetragonal PSn, Biln and Bi. lnside the complex regular structure of Bi-ySn

the mottled microstnidure consists of gray Biln and black pure Sn without any

trace of Bi inside the pSn phase. The eutectic grains consist of a binary structure

of Biln-ySn where two phases from decomposed y-Sn are present in the fom of

pSn aiong with Bi particles and the other Biln incorporated in matrix represents a

three phase decomposition. The same phenomenon was observed by Ruggiero

and Rutter (24). From this type of microstructure it is reasonable to conclude that

in the case of Bi-ySn, the Bi phase is absent due to the incorporation of Bi phase

into the Bi skeletal structure and in the case of Biln-ySn the Biln phase has

incorporated inside the Biln phase. lt can be concluded that in the present work

for statically cast samples, we have double binary structures of Bi-ySn and Biln-

ySn, segregated blocks of pure bismuth and mottled Sn dendrites. From Figure

21 it can be seen that the eutectic cell (gray, Biln-ySn binary) and complex

regular (white, Bi-ySn binary) can be distinguished from the grain boundary area.

The small complex regular structures were always observed around grain

boundaries whereas the massive complex regular structures were not confined

around grain boundaries as shown in Figure 18.

Evidence of precipitation has been observed in static cast samples inside

the Biln matrix as shown in the back-scattered SEM image in Figure 24 and also

in Figures 25 and 26 where the samples were produced at different cooling

rates. It can be observed that Bi and Sn has precipitated in the Biln matrix which

indicates that at the eutectic temperature, the Biln must have dissolved a certain

amount of Bi and Sn and precipitation occurred during cooling to room

temperature.

Figure 21 : Backscattered SEM image of eutectic cell showing segregated white

bismuth phase in the forrn of complex regular structure around grain

boundary, bismuth-indium gray spine and mottled region of tin

dendrite.

Figure 22 : EDX analysis showing the presence of three elements Bi, Sn and In

in the decornposeci structure of Bi cornplex regular structure.

c) Figure 23 : Backscattered SEM image showing a) the cubic, b) the fish spine

and c) the trigonal shaped cornplex regular structure of bismuth.

Narrow Precipitation inside matrix

tin phase

Wide ti phase

Figure 24 : Backscattered SEM image showing the precipitation of bismuth

and tin in gray rnatrix and absence of gray region where black

tin phase is narrow.

a) Cooling rate 0.2'Clmin

b) Cooling rate O.BoClmin

Figure 25 : Backscattered SEM image of complex regular structure.

a) Cooling rate 1 "Clmin

b) Cooling rate TClmin

Figure 26 : Backscattered SEM image of cornplex regular structure.

60

Figure 27 : Decomposed Sn dendrite showing several colonies or ceils having

larneilar structure.

5.2.2 Microstructure of OCC wire

The back scattered SEM images of OCC wire at a casting speed of

14mm/rnin are shown in Figure 28 a (longitudinal section) and in Figure 28 b

(transverse section). It can be seen from the longitudinal section that the overall

structure is unidirectional. There is no evidence of grain boundaries and the

segregated bismuth blocks or complex regular structure. From the transvene

section, it can be seen that al1 phases are uniformly distributed. A longitudinal

section of OCC wire produced at a casting speed of 79mmtmin is shown in

Figure 29 a and a transverse section in Figure 29 b. It can be seen from the

longitudinal section that the structure is not directional throughout the entire

length. The bottom part is not directional along the casting direction compared to

the upper part. The boundary between the upper part and the lower non-

directional part bas no segregated phase. The transverse image is identical with

that of OCC wire produced at the lower casting speed. The main difference

between OCC wire produced at different speeds is that in the case of higher

casting speed the microstructure is finer. In Figure 28 a) and b) a backacattered

SEM image of the longitudinal and transverse sections of OCC wire produced at

a casting speed of 14rnmlmin at low and high magnification shows mainly three

phases : black dendritic phase, white phase connected to black phase and a

gray matrix. An EDX analysis was carried out to check the elements present in

each phase as shown in the Figure 30. The white phase connected to black

dendrites is pure bismuth, as shown in the EDX analysis in Figure 30 a. The

black phase consists of tin rich dendrites with a trace of bismuth since the

bismuth phase is connected with tin dendrites as shown in Figure 30 b. The gray

phase in the matrix had the composition of bismuth and indium as shown in the

Figure 29 c. It can be seen that the gray matrix phase in the OCC sample has

- only bismuth and indium and no trace of tin. In statically cast sample EDX

analysis confirmed the presence of al1 the three elements, bismuth, indium and

tin in the gray matrix of eutectic ceIl, but tin was absent in the gray mottled region

inside the Bi complex regular structure which is Biln of decornposed structure.

62

There is no trace of small or massive Bi complex regular structure or Bi block

structure in the OCC microstructures but from the EDX analysis of gray rnatrix of

OCC samples it can be observed that bismuth content is higher compared to the

matrix of statically cast samples. The phases in OCC samples are uniformly

distributed and there is no trace of grain boundary at a casting speed of

14mmlmin. At a casting speed of 79mmlmin the structure is directional but one

part is parallel to the casting direction and the other part has a deviation. The

OCC wire structure produced at a casting speed of 79mmlmin was fine and it

was difficult to obtain clear images at lower magnification.

It has been observed that in case of static casting if the interphase

spacing is narrow, we may expect a structure where the binary structure of Biln-

ySn can result in the formation of a final morphology of Sn coupled with Bi in a

matrix of Biln due to the decomposition of ySn. At both casting speeds in OCC

samples it was observed that the Bi phase is coupled with Sn dendrites. It is

suggested that in the case of the OCC process at eutectic temperature the

structure was composed of ySn and matrix of bismuth and indium with higher

content of bismuth compared to statically cast samples. The decomposed

structure of y-Sn could accommodate only Bi and Sn due to the fine

microstructure. The other decomposition product, Bil n, was incorporated within

the matrix of bismuth and indium. When the Biln is incorporated in the matrix, the

morphology of the dendrite structure is different from the original and has a

similarity with the thermally annealed dendrite structure reported in the literature

(19). With OCC samples it can be concluded that the structure is not composed

of a double binary structure like statically cast samples. The ySn dendrites in the

OCC samples have decomposed resulting in the formation of PSn dendrites

coupled with bismuth and the Biln component has been incorporated within the

bismuth-indium matrix. The wire produced at a speed of 14mm/min was coarser

and an enbrged image is shown in Figure 31. The bismuth phase is not in a

complex fom but exists as blocks and is well conneded with tin dendrites,

Figure 31a and b. ln statically cast rod the precipitation of tin and bismuth was

mostly inside the gray region of the eutectic cell. EDX analysis of the gray matrix

confirmed the presence of al1 three elements, bismuth, indium and tin.

Compared with statically cast samples, the OCC samples have the

following features,

a) No segregation of bismuth in the cast samples.

b) No complex regular structure of bismuth, but rather pure bismuth connected

with tin dendrites due to the decornposition of y Sn.

C) NO evidence of double binary structures.

d) Fine and uniform distribution of al1 the phases.

Casting directio

ong itud inal section

'fan sverse section

Figure 28 : Backscattered SEM image of OCC wire, casting speed 14mmlmin.


asting directio


Figure 29 : Backscattered SEM image of OCC wire, casting speed 79rnm/min.

(cl Figure 30 : EDX spedra of OCC wire showïng,

a) White pure bismuth phase. b) Bbck dendritic tin phase,

c) Gray matrix phase without lin.

A Casting direction

a) Low magnification

b) High rnagnification

asting directio

Figure 31 : Backscattered SEM image of OCC wire, casting speed 14mmlmin,

showing a) tin dendrite and associated bismuth and b) at higher

magnification.

5.3 Compositional unifonnity in OCC samples

In statically cast products it can be seen from the SEM back scattered image

that the structure is coarse. From the structure, even at low magnification, it can

be observed that in the lower part of the cast product, the volume of segregated

white bismuth phase is more than in the upper part. The black tin dendritic

structures are distributed at random. If EDX analysis is performed at different

areas, differences in chemical composition will be observed between the upper

and lower part of the cast sample. For example, an EDX analysis of the lower

area will show significantly higher amounts of bismuth compared to the upper or

middle parts of the cast product. This indicates that for statically cast product, the

phases as well as chemical composition are not uniform.

Compared to statically cast product, OCC wire has some different features,

fine microstructure, nearly unidirectional solidification and no trace of segregation

and the phases were uniformly distributed. Although the phases were uniformly

distributed in OCC samples it was still necessary to establish the unifomity of

chemical composition at different areas of a longitudinal section. To check this,

OCC wire of 130cm in length was produced at a casting speed of 79mmhin

from one melt. The wire was cut into 13 equal pieces from start of casting to the

end of casting. Each piece or sample was polished and was prepared for EDX

analysis without etching to avoid any dissolution of a particular phase.

This is important to mention that in the present study the alloy prepared from

bismuth, indium and tin have a weight % of 57.2, 24.8 and 18 respectively.

Neutron activation analysis of a quenched sample in "Slow Poke Reador" was

conducted and to investigate a quenched sample of the alloy prepared. It was

found that the wt % of bismuth, indium and tin was 57.2, 25 and 17.8

respectively. But the result of the EDX analysis of the same quenched sample in

wt % was 51.6 % bismuth, 21.77 % indium and 26.63 % of tin. The atomic

number of indium and bismuth is 49 and 50 respectively. This irnplies that error

may occur during EDX compositional analysis of these two elements due to their

69

nearly same energy level. So the difference of actual ternary eutectic

composition and EDX result was rneasured and was added or subtracted with

the composition obtained for each element of OCC sample.

At first the weight % of bismuth, indium and tin from top and bottom areas of

the longitudinal section was measured for one sample. lt was found that in a

single sample there was no large deviation in percentage of each element from

the average weight % of individual elements measured at different locations

within the same sample.

In the next step, the purpose was to check the uniforrnity in chemical

composition, from start of casting to the end of casting and compare the

difference in wt % from the average values. The chemical composition of each

piece from different locations was measured in weight %. The average weight %

of each element was calculated from the data obtained from 13 samples. The

average reading was considered as a reference and the difference in weight %

for each element from the average value was measured as shown in Table 3. A

plot was made to describe the difference in weight % as a function of length of

sample from start of casting, Figure 32. It can be observed that the difference in

weight % from average value for bismuth, indium and tin are approxirnately 12,

11.5 and il respectively. This indicates that OCC samples have nearly uniforrn

composition.

Table 3 : Details of chemical composition of OCC samples

+O36 25.74

+1.71 25.98

+0.86 24.86

+0*38 25.73

+0.23 26.71

-1.39 26.61

-2.02 27.72

-1.71 27.23

-1.42 27.71

+0.08 27.67

+0.2 27.16

Average 26.55

In content

Difference with

average wt3C

58.14

Bi content

Avenge

Sn content

In

(wtlt)

Difference with

average wt16

Sn

(wtlb)

Difference with

average wt?h

Wt % DIFFERENCE FROM AVERAGE COMPOSITION 4 d

O & & & L O r u P C J 3 O o O

5.4 Mechanical properties

It has been observed that there are differences in structural morphology

between statically cast and OCC products. In general the statically cast products

are multicrystalline, consisting of coarse microstructure and the bismuth phase is

segregated along grain boundaries in the forrn of complex regular structure and

segregated at the bottorn due to gravit-. The matrix, which is composed of

bismuth and indium, has a lot of precipitation of bismuth and tin. In OCC, the

structure is nearly unidirectional with a tin dendrite structure in a matrix of

bismuth and indium and the phases are uniformly distributed. The bismuth phase

is closely associated with tin dendrites in the OCC products. The gray matrix has

no trace Bi phase in the form of complex regular structures or blocks. The Bi

content within the matrix is higher than in the static cast samples. Shewmon (13)

discussed the cause of embrittlement in gold alloys. He concluded that only 0.2

wt % bismuth addition in gold significantly lowered the elongation and ultimate

tensile strength of the alloy due to low solid solubility in the matrix and also the

inability to form intenetallic compounds. However in OCC samples, Bi has not

segregated and this is in accord with the high bismuth content of the matrix

cornpared to the statically cast samples.

The purpose for mechanical testing was to check the effect of rnicrostructural

change on mechanical behavior and to compare toughness as well as strength

between OCC and statically cast samples. A series of tensile tests at different

crosshead speeds were conduded using an INSTRON machine. The minimum

total length of the sample required was approximately 1OOmm.

All the statically cast samples for mechanical testing were produced with the

split mold at a cooling rate of about 0.8"CImin and the reference thennocouple

temperature was adjusted to follow a temperature profile as shown in Figure 11.

Samples from d ifferent locations of the cast rod were prepared for microstructure

observation and it was confimed that there was no difference in microstrudural

morphology. A back-scattered SEM image, Figure 33,represents the general

structure of the sample. It can be seen that bismuth blocks have segregated at 73

the bottom of the cast rod. The black phase represents tin rich dendrites. The

eutectic grains were surrounded by bismuth complex regular structure. The grey

matrix phase contains bismuth and indium.

The tensile specimens as shown in Figure 34, both from OCC and static

casting, were machined to a gauge length of 25mm. The as cast statically cast

rod products have a diameter of 3mrn and cast defects such as blow holes and

ridges were present at the surface along the entire length. To remove surface

defects it was necessary to machine out nearly imrn of thickness and the gauge

diameter was then around 2mm. In this way the segregated bismuth phase was

removed during machining. It is important to mention that the bismuth phase is

brittle in nature and the presence of segregated bismuth at the surface could

initiate cracks. Subsequent crack propagation could initiate stress concentration

at the crack tip resulting in transgranular or brittle fracture.

Surface condition is important in terms of strength, since the presence of

notch on surface increases the stress concentration at the tip of the notch and

fracture occurs at lower applied stress. In the case of OCC samples, produced at

a casting speed of 78mmlmin the surface was shiny indicating no surface

defects. However the samples were machined under identical conditions to the

static cast samples in order to have the same surface appearance and condition

after machining.

It was observed that during machining the surface rnaterial of statically

cast samples came out in the form of chips or small particles and for OCC

samples the surface material was in the shape of continuous turnings as shown

in Figure 35. The small particle or chip shape is a good indication of brittleness

and the continuous turnings indicate ductility. Again the static samples fractured

if the depth of cutting exceeded 0.1 5mm per pass. This was not the case with

OCC wire. The shape of the chips and limitation of the depth of cut indicates the

low energy of statically cast products compared to OCC wire. Mechanical testing

was carried out for both staüc and OCC samples at different crosshead speeds

to check the extent of strain sensitivity. The crosshead speed was selected at

i.25mmlmin. 2.5mmlmin, 5mmhnin, 7.5rnmlmin and 1 Ommlmin and kept

constant until fracture occurred. The initial strain rate was calculated from the

crosshead speed. Frorn this test a plot of stress vs. elongation % was obtained at

different crosshead speeds. The crosshead speed and the corresponding initial

strain rate have been tabulated in Table 4. The terni initial strain rate is the ratio

of crosshead displacement in mmlsecond and initial gauge length in mm.

Table 4 : Relation between crosshead speed and initial strain rate

CROSS-HEAD SPEED (MMIMIN) 1 INITIAL STRAIN RATE ( SEC*')

Figure 33 : Backscattered SEM image of statically cast rod accepted

for tensile test. Segregated bismuth was removed by machining

during sample preparation.

Gauge diameter I

Gauge length = 25 mm

Figure 34 : Schematic diagram of tensile specimen

76

OCC

Static

Figure 35 : Photograph of tumings to show difference in material properties between OCC and statically cast samples.

The tensile tests were conducted at room temperature (25°C). For a particular

crosshead speed the tensile test was conducted for three OCC samples and

three statically cast samples. At a particular crosshead speed for both types of

samples maximum and minimum values of stress and elongation % were

observed. Comparative plots of stress vs. elongation % at different crosshead

speeds were prepared for each sample, Figure 36 to Figure 40.

At crosshead speeds from 1.25mm/min to 10mmlmin the OCC samples

have an elongation value within the range 3560% of initial gauge length. In the

statically cast products the range was 2243% at crosshead speeds from

1.25mmlmin to Smmlmin and the maximum and minimum elongation observed at

a crosshead speed of 7.5mmlmin was approximately 20% and 1% respectively.

At the maximum crosshead speed of 10mmlmin the statically cast product had

no trace of plastic deformation. This indicates that statically cast samples had

inconsistency in terms of elongation or ductility. The consistent range of

elongation at different crosshead speeds for OCC samples indicates that this

material is not strain rate sensitive in contrast to static cast samples.

The values of yield stress (0.2% off set), UTS and maximum elongation for

each set of samples at different crosshead speeds were determined and the

details for each sample are shown in Table 5 for statically cast samples and in

Table 6 for OCC samples.

OCC Static

a) Maximum elongation values

OCC Static

b) Minimum elongation values

Figure 36 : Comparison of elongation values between OCC and statically cast samples at a crosshead speed of 1 -25mmlmin.

OCC


-

OCC


Static

Static

Figure 37 : ComparÎson of elongation values between OCC and statically cast samples at a crosshead speed of 2.5mmlmin.

OCC


OCC


Static

Figure 38 : Cornparison of elongation values betwean OCC and statically cast samples at a crosshead speed of Smmlmin.

OCC Static


OCC


Static

Figure 39 : Cornparison of elongation values between OCC and statically cast samples at a crosshead speed of 7Smmlrnin.

O 10 20 30 40 50 ôû

Elongation %

OCC


O 10 20 30 40 50 60

Elongation %

OCC


Static

O 10 20 30 40 50 60

Elongation %

Figure 46 : Cornparison of elongation values between OCC and statically cast samples at a crosshead speed of IOmrnfmin.

Table 5: Mechanical properües of statically cast samples

" - Premature failure. CHS - Crosshead speed, RA - Reduction in area after

fracture, YS - Yield stress, UTS - Ultimate tensile strength

Table 6: Mechanical properties of OCC samples

1 Sample ( CHS (mmlmin) 1 YS (MPa) 1 UTS (MP~)/ RA % 1 ELONGATION %

CHS - Crosshead speed, RA - Reduction in area after ftacture, YS - Yield

stress, UTS - Ultimate tensile strength

A distribution of yield stress, ultimate tensile strength and elongation with

increasing crosshead speed is show in Figures 41,42 and 43 respectively.

The yield stress is the stress at which a specific measurable plastic

deformation, cornmonly 0.2 % plastic strain has taken place. The yield stress of a

given alloy depends on the microstructure, test temperature and strain rate. In

the present work the test temperature and strain rate in ternis of crosshead

displacernent were kept constant for both OCC and statically cast samples.

From the plot of yield stress vs. crosshead speed shown in Figure 41, it

can be observed that for OCC samples the values of yield stress at all crosshead

speeds are consistent in contrast to the behavior observed with specirnens which

had been statically cast. In the case of statically cast product large differences

between maximum and minimum yield stress values at crosshead speeds of

7.5mmlmin and lOmmlmin can be observed. The minimum value of yield stress

has a large deviation from the average value. This indicates that the samples had

premature failure either without, or with only a small amount of plastic strain. The

small areas under the stress-strain curves clearly indicate the low energy for

fracture that is consistent with brittle behavior and indicates the low toughness of

the statically cast samples. In OCC samples the deviation between maximum

and minimum yield stress at al1 crosshead speeds is less and there is an almost

linear relationship between yield stress and crosshead speed. These results

indicate that even at high crosshead speed beyond Srnmimin, OCC samples

have no indication of premature failure.

In the case of OCC samples there exists no deaease in yield stress but

in statically cast samples the yield stress drastically decreases with the increase

in crosshead speed due to premature failure. Beyond a crosshead speed of

Smmlmin, most of the samples showed a premature failure, which indicates that

the statically cast samples, are unable to withstand the stress beyond a

crosshead spaed of Smm/min. This implies that with the test temperature and

constant strain rate the reason for higher toughness observed with OCC samples

is the result of fine microstructure and uniform distribution of al1 the phases.

In Figure 42 it can be seen that the statically cast products not only have

a large difference between maximum and minimum values of ultimate tensiie

strength, but the UTS also decreases as the crosshead speed increases. This

implies that the lower values of UTS for statically cast samples at high strain

rates are not tnie UTS but rather correspond to maximum stress before

catastrophic failure which is very low compared to OCC samples. In the case of

OCC samples the deviation is much less compared to the static samples and

there is no evidence of premature failure. It is evident that the statically cast

sarnples are strain sensitive in contrast to the OCC samples.

In Figure 43 the plot of elongation vs. crosshead speed shows that OCC

samples have a higher elongation compared to static samples. At higher

crosshead speeds the statically cast products have very low elongation and at

10mmlmin there is practically no elongation due to premature failure. High

elongation indicates high ductility and in this context OCC samples have superior

characteristic compared to statically cast samples.

It can be seen from these graphs that in terms of yield stress and ultimate

tensile strength, OCC products have higher values over a wide range of strain

rates. The elongation of OCC wire in general has a higher range of values at low

to high strain rates. This confimis that OCC material has a higher ductility and

toughness than statically cast products due to the improved structural

morp holog y.

O 1.2s 2 5 3.76 6 6.26 7.6 0.76 10 11.2s 125

CROSSHEM SPEED ( mm l min ) I OCC - - - - - STATlCALLY CAST SAMPLE

Figure 41 : Plot of yield stress vs. crosshead speed showing the low value of yield stress at higher aosshead speeds which produces prernature failure of the staticaîly cast sarnples and aie high and consistent yield stress values of OCC samples.

Figure 42 : Plot of ultimate tensile stress vs. ciosshead speed showing, a) significant inconsistency in UTS for statically cast samples at higher crosshead speed and b) higher and more consistent UTS of OCC sarnples.

OCC

Figure 43 : Plot of elongation vs. crosshead speed sttowing the greater duaility of OCC samples cornpared to statically cast samples.

5.5 Fracture surface of OCC and statically cast samples

It has been observed from mechanical testing and the data derived in

terms of elongation, yield stress and ultimate tensile stress that the OCC

samples in com parison with statically cast samples have in al1 respects higher

values. At al1 crosshead speeds the OCC samples absorbed a higher amount of

energy in the plastic range, which is the area under the stress-strain curve

indicating the higher toughness compared to statically cast samples. In the case

of statically cast samples, the energy absorbed was significant at lower

crosshead speed, but at higher crosshead speed the energy absorbed was

practically ni1 and failure was catastrophic. To investigate the fracture surface,

both OCC and statically cast samples were identified according to crosshead

speed and examined using SEM secondary images.

Figure 44 and 45 show the fracture surfaces of OCC and statically cast

samples for crosshead speeds of 1.25mmlmin and 2.5mmlmin respectively. The

fracture surfaces are ductile in nature and there is no trace of cleavage. Both

OCC and static samples exhibit necking effects and from the fracture surface it

can be determined that the reduction in area is approximately 95%.

In Figure 46 the fracture surface for both OCC and statically cast samples

tested at a crosshead speed of Smmlmin are essentially similar. Necking and

high reduction in area for both types of samples are clearly evident.

A different fracture surface is obtained with crosshead speeds of

7.5mmîmin and 10mm/min. As shown in Figure 47 a, the OCC samples,

fractured at a crosshead speed of 7.5mm/min, have a large reduction in area and

the fracture surface shows no trace of cleavage. In the case of the statically cast

product in Figure 47 b, there is evidence of plastic deformation and a slight

necking effect, but the fracture surface shows distinct cleavages with no trace of

dimple rupture. This indicates that at this strain rate, the statically cast product is

brittle in nature.

Again at a crosshead speed of 10mmImin it can be seen from Figure 48

a, that the OCC sample has a high reduction in area and the fracture surface

appearance is dirnpled. The statically cast product, Figure 48 b, fractured at the

same crosshead speed, has no necking effect and this is consistent with the

observed catastrophic failure.

In Figure 49 a and b the fracture surfaces at higher magnification are

shown for both OCC and statically cast samples respectively, fractured at a

crosshead speed of 1.25mmlmin. It can be seen from Figure 49 a, that the OCC

samples have a dimpled appearance. Another feature is that in OCC samples the

micro voids have almost rounded dimples. In the statically cast samples

rectangular shaped brittle particles are present. Void formation begins at the

matrix- particle interface due to decohesion or fracture of the brittle particle inside

the matrix. Subsequently the voids coalesce to cause final fracture. It is

suggested that void formation in the case of OCC is due to decohesion of

precipitated round shaped particles from the matrix and Bi particles attached to

Sn dendrites. In the case of statically cast samples, void formation is due to

fractured particles segregated along grain boundaries and to decohesion of the

precipitated phase in the matrix.

At a crosshead speed of lUmm/min, the OCC sample, Figure 50 a, has a

dimpled rupture surface and from the high elongation this can be considered as

ductile fracture. The statically cast sample, Figure 50 b, shows a fracture surface

which consists of cleavages and this is typical of brittle fracture. This is consistent

with the stress vs. elongatîon behavior where the energy absorbed was low and

the product behaved in a brittle rnanner. Another feature from Figure 50 b for the

statically cast sample is that crack propagation is along the grain boundaries.

To investigate the reason for premature failure in the statically cast

products a sample was polished and examined before the tensile test. It was

found as shown in Figure M that there are voids along the grain boundaries in

the cast product as indicated by the arrow. These voids cause stress

concentration and enhance crack propagation along the grain boundaries and

lead subsequently to transgranular fracture at high crosshead speeds.

a) OCC sample

b) Statically cast sample

Figure 44 : Fractography after tensile test at a crosshead speed of 1.25mm/min.

60th OCC and statically cast samples are dudile in nature showing

high reducîion in area.

a) OCC sample


Figure 45 : Fractography after tensile test at a crosshead speed of

2Srnmlmin showing no significant difference in reduction

of cross-sectional area-

a) OCC sample


Figure 46 : Fractography after tensile test at a crosshead speed of Smmlmin

showing the evidence of ductility for both OCC and statically

mst samples.

a) OCC sample

b) Staticall y cast sample

Figure 47 : Fractography after tensile test at a crosshead speed of 7.5mm/rnin

showing the dudile fracture of OCC sample and brittle fracture

of statically cast sample.

a) OCC sample


Figure 48 : Fractography after tensile test at a crosshead speed of

1 Omrnimin showing, a) high reduction in cross-sectional area

and ductile nature of fracture surface of OCC sample,

b) cleavages and no reduction in area on fracture surface of

statically cast sarnple.

a) Fracture surface of OCC sample showing srnall microvoids and dimples

b) Fracture surface of statically cast sample showing large microvoids

with dimples and cleavage fracture of a particle as shown by arrow

sig n

Figure 49 : Fractography of tensile test sample at a crosshead speed of

1.25mmlmin and at high magnification.

a) Fracture surface of OCC sample which is ductile in nature

b) Fracture surface of statically cast sample showing cleavages and

evidence of intergranular fracture as shown by anow sign

Figure 50 : Fradography of tensile test sarnples, at a crosshead speed

of 1 Ommlmin and at high rnagnification.

Figure 51 : SEM secondary image showing voids along the grain boundary

of a statically cast sample.

Crack nucleation mechanisms Vary according to the type of material which

may be brittle, semi-brittle or ductile. In the statically cast samples, the bismuth

content is concentrated around the grain boundaries. In contrast, OCC samples

exhibited no trace of segregated bismuth. At a cast speed of 79mm/min, even

with the existence of more than one crystal there is no segregated bismuth phase

or voids around the grain boundaries. Since bismuth is brittle in nature, the

dislocations are practically immobile, whereas in ductile materials there is

relatively little restriction on dislocation movernent. Again material heterogeneity

can produce a stress concentration which can nucleate a crack. With respect to

the surface condition of the OCC and statically cast samples, since al1 the

specimens were machined under the same conditions it is assumed that stress

concentration due to surface condition which in this case may be steps, striations

or depressions, is the same. In the interior of the statically cast samples there

exist voids and segregated brittle bismuth phase along grain boundaries.

As mentioned previously both W C and statically cast products behave in

a ductile manner up to a crosshead speed of Smmfmin. However at higher

crosshead speeds, OCC samples remain ductile, white the statically cast

samples show brittle behavior. For OCC samples, the fracture surface is dimpled,

and the energy required for complete fracture is higher than that for the statically

cast product, at all the crosshead speeds.

At crosshead speeds beyond 5mmlmin statically cast samples exhibit the

features of cleavage fracture and low strength prior to fracture which are typical

characteristics of brittle fracture behavior, thus there is a ductile to brittle

transition with increase in strain rate.

CHAPTER 6

CONCLUSIONS AND FUTURE WORK

6.1 Conclusions

From the results of this experimental study the following conclusions can be

drawn,

1. A ternary eutectic alloy in the Bi-ln-Sn system which is brittle in nature due to

its high bismuth content (57.2 wt %) and difiicult to produce by conventional

methods, was successfully processed in wire form with small cross-sectional

area by the "Ohno Continuous Casting" process. The net shape product had

a high quality surface finish. In contrast the statically cast alloy exhibited

surface defects such as ridges and blowholes.

2. The OCC wires had fine microstructure and al1 the phases were distributed

unifomly without any trace of segregated bismuth. In statically cast samples,

bismuth complex regular structures were observed along the grain

boundaries of eutectic cells. Bismuth blocks and massive bismuth cornplex

regular structures were prone to gravity segregation in the temperature

plateau region of statically cast samples and this phenornenon yields non-

uniformity in phase distribution.

3. Statically cast samples exhibited regions of double binary structure similar to

that observed by Ruggiero and Rutter. One binary structure is Biln-ySn and

the other is Bi-ySn. However in OCC samples, double binary structures were

absent and the structure was composed of ySn and bismuth-indium matrix.

The bismuth content of the bismuth-indium matrix in OCC samples was

higher than that in statically cast samples.

4. Several colonies or cells having lamellar structure were observed in a slowly

cooled, coarse, tin-rich dendritic structure. The growth of each colony or cell

indicates an individual nucleaüon event and it can be concluded that the ?Sn

dendrites have decomposed below the eutectic temperature by an eutectoid

reaction . 5. To obtain accurate tensile test data with materiai which is strain sensitive or

ductile in nature, a modified grip was designed for use with the INSTRON

machine. This not only prevented plastic deformation or crack of the

specimen around grip area due to the high impact and compressive force of

INSTRON grip but also eliminated slip during loading.

6. A significant improvernent in mechanical properties was achieved with

samples produced by the OCC process. It was found that OCC samples are

ductile even at a crosshead speed of lOmm/min, whereas statically cast

samples had premature failure at a crosshead speed above 5mmhin.

7. It has been established that a lead free, low melting point alloy can be cast

continuously in the fonn of wire even with material which is normally brittle in

nature and the mechanical properties of the products are superior compared

to statically cast samples. Hence it would be useful, to study some additional

bismuth bearing, lead-free solder alloys which could be industrially acceptable

but would be difficult to produce by conventional methods.

6.2 Future work

1. It was found that the OCC samples have a higher bismuth content in the

matrix compared to statically cast samples. It was also found that the bismuth

complex phases were absent in OCC samples. This suggests that the matrix

of OCC samples is no longer a composition of Biln, but may consist of

metastable phases such as Bi& or BMn. In future work, the composition of

the matrix of OCC samples should be investigated in more detail in order to

clarify this aspect.

It was observed that the microstructure of OCC samples produced at a

casting speed of 79 mmlmin was not directional, in contrast to samples

produced at a casting speed of 14 mmlmin. Also it was found that in the case

of statically cast samples, the structure was directional within the temperature

gradient region but there exists a bismuth complex regular structure. This

might also exist in OCC samples produced at a very slow casting speed.

Hence it is important to determine the optimum conditions for the OCC

process in order to investigate the effed of temperature gradient and cooling

rate at which the bismuth complex regular structure will be replaced by the

formation of a matrix with a high bismuth content and no bismuth segregation.

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Documents

COMPARATIVE STUDY Bi-lnSn TERNARY - TSpace … · COMPARATIVE STUDY OF Bi-InSn ... 2mm diameter by the Ohno Continuous Casting (OCC ... between the mold wall and the cast strand creates