RAPID SOLIDIFICATION PROCESSING OF INDIUM GALLIUM ANTIMONIDE ALLOYSSecure Site  · 2020. 4. 2. ·...

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RAPID SOLIDIFICATION PROCESSING OF INDIUM GALLIUM ANTIMONIDE ALLOYS Item Type text; Thesis-Reproduction (electronic) Authors Kumta, Prashant Nagesh, 1960- Publisher The University of Arizona. Rights Copyright © is held by the author. Digital access to this material is made possible by the University Libraries, University of Arizona. Further transmission, reproduction or presentation (such as public display or performance) of protected items is prohibited except with permission of the author. Download date 17/04/2021 03:54:05 Link to Item http://hdl.handle.net/10150/276468

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RAPID SOLIDIFICATION PROCESSING OFINDIUM GALLIUM ANTIMONIDE ALLOYS

Item Type text; Thesis-Reproduction (electronic)

Authors Kumta, Prashant Nagesh, 1960-

Publisher The University of Arizona.

Rights Copyright © is held by the author. Digital access to this materialis made possible by the University Libraries, University of Arizona.Further transmission, reproduction or presentation (such aspublic display or performance) of protected items is prohibitedexcept with permission of the author.

Download date 17/04/2021 03:54:05

Link to Item http://hdl.handle.net/10150/276468

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Order Number 1331410

Rapid solidification processing of indium gallium antimonide alloys

Kumta, Prashant Nagesh, M.S.

The University of Arizona, 1987

U M I 300 N. ZeebRd. Ann Arbor, MI 48106

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RAPID SOLIDIFICATION PROCESSING OF

INDIUM GALLIUM ANTIMONIDE ALLOYS

by

Prashant Kumta

A Thesis Submitted to the Faculty of the

DEPARTMENT OF MATERIALS SCIENCE AND ENGINEERING

In Partial Fulfillment of the Requirements For the Degree of

MASTER OF SCIENCE

In the Graduate College

THE UNIVERSITY OF ARIZONA

19 8 7

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STATEMENT BY AUTHOR

This thesis has been submitted in partial fulfillment of requirements for an advanced degree at The University of Arizona and is deposited in the Uhiversity Library to be made available to borrowers under rules of the Library.

Brief quotations from this thesis are allowable without special permission, provided that accurate acknowledgement of source is made. Requests for permission for extended quotation from or reproduction of this manuscript in whole or in part may be granted by the head of the major department of the Dean of the Graduate College when in his or her judgement the proposed use of the material is in the interests of scholarship. In all other instances, however, permission must be obtained from the author.

SIGNED:

APPROVAL BY THESIS DIRECTOR

Tl is has been approved on the date shown below:

- h • 7- /£> — 8 ~7 S. H. Risbud

Professor of Materials Science and Engineering

Date

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To my dearest Parents "Nagesh and Sumati Kumta"

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ACKNOWLEDGEMENTS

I wish to thank my advisor, Dr. Subhash Rlsbud for the constant

encouragement, advice, and assistance needed to complete this study. I

also wish to express my gratitude to Dr. David Poirier for his valuable

suggestions during the writing of this thesis. Sincere thanks are also

extended to Mr. Dick Van Reeth for his help in machining the melt

spinning set up needed for this study. I am also grateful to my dear

friend Mr. V. P. Dravid who performed electron microscopy on the

rapidly solidified flakes with the microscopy facility at Lehigh

University. The help of Mr. D. S. Mishra of the Pharmaceutical Science

Department at the University of Arizona who performed the DSC analysis

is also appreciated. Finally, I would like to thank my parents, who are

mainly responsible for what I have accomplished.

ill

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TABLE OF CONTENTS

Page

LIST OF ILLUSTRATIONS v

LIST OF TABLES x

ABSTRACT xi

CHAPTER

1. INTRODUCTION 1

2. BACKGROUND AND LITERATURE 4

2.1 Rapid Solidification Technology 4 2.2 III-V Semiconductors 21 2.3 Ternary III-V Antimonides 21 2.4 III-V Ternary phase diagrams 26 2.5 Synthesis Of Bulk Crystals Of

InxGa1_xSb 28

2.6 Thin Film Methods 33

3. EXPERIMENTAL SET UP AND PROCEDURE 35

3.1 Materials Preparation 35 3.2 Materials Characterization 40

4. RESULTS AND DISCUSSION 44

4.1 X-ray Diffraction Analysis 44 4.2 Surface Morphology 53 4.3 Thickness Measurements 63 4.4 Microstructure Analysis 67 4.5 Transformations 85 4.6 Electrical Property Measurements 93

5. SUMMARY AND CONCLUSIONS 97

REFERENCES 99

iv

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LIST OF ILLUSTRATIONS

Figure Page

1. Cooling rates and various structural forms achieved in rapid quenching from the liquid state 5

2. Schematic for representing the concept associated with undercooling of the a phase for a simple eutectic system 7

3. Extended phase concept for adiabatic solidification of the a phase 9

4. Schematic concept for undercooling development of metastabler phase and a special case of hyper cooling where 7 phase was suppressed leading to formation of an amorphous phase 11

5. Microstructural consequences of rapid solidification 12

6. Schematic representation of various techniques of RSP 17

7. Various techniques of filament casting 19

8. Powder preparation techniques 22

9. The InSb-GaSb pseudobinary phase diagram 29

10. Schematic diagram of synthesis of InxGa1_xSb alloys 36

11. Schematic of the RSP set up 38

12. X-ray trace of ingot, flake and powder obtained from melt of run 19 50

v

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vi

LIST OF ILLUSTRATIONS — Continued

Figure Page

13. (a) SEM micrograph showing surface of flake obtained from run 9 54

(b) SEM micrograph showing surface of flake obtained fran run 20 54

(c) SEM micrograph showing precipitates in flakes obtained form run 1 55

14. ED AX analysis of the precipitate observed on the flake surface 57

15. EDAX analysis of the matrix on the flake surface . 58

16. (a) SEM micrograph of powder obtained from run 10 59

(b) SEM micrograph of powder obtained from run 10 59

(c) SEM micrograph of powder obtained from run 10 60

(d) SEM micrograph of powder obtained frcxn run 10 60

(e) SEM micrograph shewing fractured powder particle obtained fran run 10 61

(f) SEM micrograph shewing powder obtained from run 9 61

(g) SEM micrograph showing powder obtained from run 8 62

17. (a) SEM micrograph of cross section of flake obtained from run 1 64

(b) SEM micrograph of cross section of flake obtained from run 19 64

18. Variation of flake thickness with disk velocity 65

19. Plot of experimentally determined flake thickness (d, microns) versus cooling rate ( K/s) calculated assuming validity of the equation Cd =81 68

4

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LIST OF ILLUSTRATIONS — Continued

Figure Page

20. (a) Optical micrograph showing cross section of ingot of composition IriQ 46GaQ 54Sb 70

(b) Optical micrograph showing the etched surface of the flake obtained fran run 7 70

(c) SEM micrograph shewing the etched surface of the flake observed in Figure 20(b) 71

21. (a) Optical micrograph showing the cross section of flake obtained fran run 7 73

(b) Optical micrograph showing the cross section of flake obtained from run 14 73

(c) SEM micrograph showing the cross section of flake observed in Figure 21(b) 74

22. (a) SEM micrograph showing the etched cross section of powder obtained from run 14 75

(b) Optical micrograph showing the etched cross section of powder obtained from run 14 75

(c) Optical micrograph showing the etched cross section of pcwder obtained from run 10 76

23. (a) Optical micrograph showing the etched surface of the flake obtained frcm run 18 78

(b) Optical mocrograph shewing etched surface of the flake obtained from run 20 78

(c) optical micrograph showing the etched surface of the flake obtained from run 19 79

24. SEM micrograph of etched surface of flake observed in Figure 23(a) 79

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25

26

27

28

29

30

31

32

33

34

35

36

LIST OF ILLUSTRATIONS — Continued

Variation of grain size with disk velocity

Bright field image of flake obtained from run 13

Diffraction pattern obtained frcm an area of flake observed in Figure 26

Diffraction pattern obtained for powder obtained from run 20

Bright field image of the same region from which diffraction pattern in Figure 28 was obtained

Diffraction pattern obtained for powder obtained from run 20 after exposing the sample to the electron beam for 10 minutes .

Dark field image of the region from which diffraction pattern in Figure 28 was obtained

Auger electron spectra showing presence of C, 0, Cu, In, Ga, Sb in flake obtained frcm run 16

Depth profile for Cu, C and elemental 0 in flake obtained from run 16

DSC traces of ingot and flakes obtained from run 18, 19 and 20

DSC traces of corresponding powders obtained from run 18, 19 and 20

DSC traces of powder obtained from run 8, 9 and 10

viii

Page

80

82

82

83

83

84

84

86

87

88

90

92

1

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ix

LIST OF ILLUSTRATIONS — Continued

Figure Page

37. Variation of bulk electrical resistivity with annealing temperatures for flakes obtained from run 7 94

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LIST OF TABLES

Table Page

1. Range of dendritic spacings associated with cooling rate and processes 14

2. Sunmary of cooling rate measurements for splat cooling 15

3. Quench rates and sample shapes in splat quenching 18

4. Fiber processing techniques 20

5. Powder processing techniques 23

6. Selected properties of III-V compounds, germanium and silicon 24

7. Possible combinations of the III-V antimonides 25

8. Lattice parameters of binary constituents in III-V ternary antimonides 27

9. Melt spinning compositions with melt spinning parameters 45

10. Lattice parameter values 49

11. Summary of XRD and DSC analysis 91

12. Resistivity values cited in literature 96

x

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ABSTRACT

Solidification from the melt is an essential step in nearly all

conventional processes to produce bulk materials for industrial

applications. Rapid quenching from the liquid state at cooling rates of

102 to 106K/S or higher has developed into a new technology for

processing novel materials. I^Ga^.^b a ternary III-V compound

semiconductor was synthesized by using the rapid spinning cup (RSC)

technique. Several compositions of these alloys were batched and cast

into ingots in evacuated sealed quartz tubes. These ingots were then

melted and ejected onto a rapidly rotating copper disk. This resulted

in the generation of flakes or powders depending on the rpm of the

disk. Mlcrostructural characterizatlcn of the flakes and powders was

performed using XRD, SEM and TEM. Efforts were also made to measure the

bulk resistivity of the annealed flakes to see the effect of annealing

on ordering of the phases.

xi

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CHAPTER 1

INTRODUCTION

The Increasing Interest In the field of microelectronics has

led to a vigorous and rapid development of semiconductor materials

science, technology and device fabrication. The last several years have

witnessed the large scale application of silicon in new devices with

enormous size reductions and packing densities. The real emergence of

silicon as a semiconductor material Is based cn Its unique advantages,

although It possesses certain limitations. The band gap of silicon and

germanium are indirect and are fixed at definite values of Eg(Ge) =

0.68 eV and Eg(Si) * 1.1 eV. Owing to these factors they are not useful

as semiconductor light sources or microwave devices; the energy band

structure Is suitable neither for light emitting diodes and lasers, nor

for transferred electron devices ( Gunn diodes ). This is the reason

why the interest in compound semiconductors is growing continuously

with an examination of their capability for electronic applications[ 1 ].

Among the possible compound semiconductors the III-V compounds

are the most Important ones. The group consists of compounds which are t

composed of the group IIIA and the group VA elements, however the

compounds containing boron and nitrogen are generally omitted primarily

due to their physical and chemical nature. The growth of large single

1

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crystals, epitaxial layers or p-n junctions of these materials is very

difficult. Combination of these six remaining elements ( P,As and Sb in

the group V and Al, Ga and In in the group III ) gives nine binaries,

eighteen ternaries and fifteen quaternary compounds. Only parts of

these combinations are known yet, and among them one can find compounds

as GaAs, GaSb, InSb, GaP, GaAsP, InGaAs, GaAlAs or GalnAsP which have

great technical importance forming the material base of industrial mass

production (LED's, GaAs devices etc ) and optoelectronic and microwave

devices.

The III-V compounds are used in optoelectronics and microwave

technologies. Furthermore new applications of these compounds such as

high speed digital integrated circuits (switching rates of less than

lOOps at very low power level) are also promising. At present a large

amount of data for III-V binaries are available but interest has

increased recently in some mixed III-V alloy systems. Compared to the

continuously growing effort in developing appropriate technological

methods, for III-V binary compounds, the preparation techniques for

ternary and quaternary III-V compounds have yet to fully exploit the

emerging rapid solidification technology currently used mainly for

metallic systems.

The present work is thus focussed on rapid solidification

processing (RSP) of Ir Ga ^ Sb alloys, of interest in galvanomagnetic

devices and as infra red detectors in the spectral range 1.8 - 6.9 fim.

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The major objectives of the research were synthesis and

characterization of rapidly solidified flakes and powders using X-ray

diffraction (XRD), scanning electron microscopy (SEM), transmission

electron microscopy (TEM) and Auger electron spectroscopy (AES)

techniques.

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CHAPTER 2

BACKGROUND AND LITERATURE

2.1 Solidification Technology

The technique of quenching or rapid cooling in the solid state

has been practiced fay metallurgists and material scientists [2-5] in

order to develop suitable microstructures and obtain desired

properties. However quenching from the liquid state in order to modify

the process of solidification to obtain a solid has in the recent years

developed into a new technology known as Rapid Solidification

Technology (RST).

The cooling rates and the corresponding various structural

forms achieved In rapid solidification is Indicated in Fig.l. Starting

from the liquid state [6], a cooling rate of 106K/s or higher can yield

metallic glasses. Cooling rates in the range from 102 to 106K/s produce

crystalline alloys, which are called, depending an their average size,

micro or nano crystalline. Hie rapidly solidified materials have the

form of thin ribbons, wires, flakes or powders.

It is currently recognized that rapid solidification affords i

numerous opportunities in materials development.[7]. The primary

virtues afforded by rapid solidification are a) Chemical homogeneity,

or decreased chemical segregation, b) refined grains and other

4

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LIQUID STATE |

METALLIC 10* GLASSES

.10* NANO MICRO

QUENCHING CONVENTIONAL METALLURGY

.9* -

COOUNG RATE (Ks" I

Figure 1. Cooling rates and various structural fox'ms achieved in rapid quenching from the liquid state.[6]

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structures, c) increased solute solubility above conventional

equilibrium levels and d) metastable phase development which includes

amorphous and microcrystalline structures.

The concept of chemical segregation development during rapid

solidification of an alloy accompanied by extensive undercooling can be

visualized from a simple eutectic phase diagram shown in Fig.2. This

model illustrates the concept of undercooling, where nucleation of the

solid phase sets in at temperatures significantly below the liquidus

temperature (Tj) of the alloy at hand. Two assumptions are applied to

the model a) heterogeneous nucleation sites are minimal and b)

adiabatic conditions, i.e the latent heat of solidification AH is

absorbed fast enough within the system so that, effectively, any heat

removal from the system during solidification is inconsequential. The

model shown in Fig.2 illustrates three undercooling conditions from the

liquid state and is associated with temperatures T1( T2 and T3. The

degree of undercooling is associated with the recalescence temperature

Tr (temperature obtained from the release of latent heat of

solidification), and the solidus temperature Tg.

Case l.[7]: hypocooling where undercooling from the liquid state

to Ta occurs but the latent heat of solidification raises the

temperature above the solidus temperature, i.e TR > Tg. In this

instance, local equilibrium and solute partitioning occurs since Tr is

in the a + L field. This condition leads to dendritic growth with

solute rich cores, i.e. chemical segregation. The dendrites become

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Of + B

V.B

Increasing undercooling

Critical Hypocooling undercooling Hypercooling

Homogeneous » Segregated Homogeneous a

Time Time Time

Figure 2. Schematic for representing the concept associated with undercooling of the a-phase for a simple eutectic system. [7]

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8

larger as TR increases.

Case 2T71: Undercooling from the liquid state to T2 and in

this instance TR = Tg. This condition is described as critical

undercooling and is associated with

AH =T2/TsCp1 dT (1)

where A H is the latent heat of solidification, and Cp1 is heat

capacity of the liquid phase. In this case solidification can occur

massively within the a phase without compositional piartitioning and

local equilibrium at the solid/liquid interface.

Case 3[7]: hypercooling and is represented by TR < Tg and,

similar to the previous condition, results in complete solute trapping

hence segregationless. solidification. As a result the composition of

the a phase is uniform and is equal to CQ. The last two conditions

feature very rapid nucleation, and as a result fine or microcrystalline

grains are expected to dominate the morphology.

The same concept of undercooling and recalescence can be

extended to the understanding of the formation of metastable phases.

This concept induced by rapid solidification is shown in Fig. 3 where

once again a simple binary eutectic is used for the illustration.

Again, adiabatic solidification is assumed. If the melt can be

sufficiently undercooled, the liquidus and solidus boundaries can be

extended such that the two phase region of « + L in Fig.3 will coexist

at temperatures below the equilibrium eutectic leading to a condition

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T1

%B »-

Figure 3 . Extended phase concept for solidification of thea-phase.[7]

adiabatic

<0

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where the /3 phase is totally suppressed. In principle conditions of

hypocooling, critical undercooling and hypercooling can be generated

(see Fig.2) producing a range of extended a phase morphologies i.e.

dendritic (segregation), cellular, and completely homogeneous. The

degree to which phases, such as a shown .in Fig.3 can be extended using

current rapid solidification techniques, is probably not large but

significant.[7,8]

Other forms of metastable phase development afforded by rapid

solidification can lead to new and unique phases. An example of such

behaviour is shown in Fig. 4, where the metastable r phase forms at the

expense of the more stable a and 0 phases as a result of appropriate

conditions of undercooling. If a condition of critical undercooling

occurs, the metastable y phase can form in a homogeneous manner. Also,

it is conceivable that if the amount of undercooling, i.e. below the

glass transition temperature Tg, is very large then the heat of

solidification, AH, is zero, and the metastable r phase can be

suppressed. This would lead to the formulation of an amorphous phase.

Microstructural consequences of rapid cooling from the molten

state are summarized in Fig.5. In passing from ordinary casting

practice to cooling rates greater than 102 K/s, the microstructural

features become refined because the time for coarsening during

solidification is reduced. [8]

The morphological assessment of rapidly solidified materials

has been mainly associated with dendritic behaviour. Jones[9],

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Increasing undercooling

Figure 4. Schematic concept for undercooling development of metastable 7 phase and a special case of hyper coo ling where 7 phase was suppressed leading to formation of an amorphous phase.[7]

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rr Coarse dandritas, eutectlcs and other mlcroconstltuants

Incraasing cooling rata (Kfs)

Refined microstructures

Composition and process Fine dendrites,

eutectlcs and other mlcroconstltuents

Refined microstructures dependent r

Fine dendrites, eutectlcs and other mlcroconstltuents

4J-Novel microstructuraa

Extended solid solutions, mlcrocrystalllne structures, metastable crystalline phases, amorphoua solids

Novel microstructuraa

Extended solid solutions, mlcrocrystalllne structures, metastable crystalline phases, amorphoua solids

Increasing homoganalty

Figure 5. Mlcrostructural consequences of rapid solidification.[8]

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Mehrabian[10], Grantll] and Flinn[7] have provided a correlation for

dendritic spacings provided by solidification from the melt as a

function of cooling rate. This correlation is shown in Table 1. A

detailed discussion on the determination of the cooling rate from the

microstructures obtained by rapid solidification can be found in the

book written by Flinn[7]. However, it has also been possible to

estimate the cooling rate from the specimen thickness. Jones[12] has

discussed in detail the various types of cooling conditions that can be

obtained and the corresponding agreement of the theoretical and

experimental values of the cooling rate. The available experimental

measurements of cooling rate by Predecki et al.[13], Harbur et al.

[14], Miroschnichenko and Zakharov [15], Lohberg and Muller [16,17],

Brower et al. [18] are summarized in Table 2. They range from about 105

- 10® K/s for splat thicknesses ranging from about 100/zm down to 1 jum.

Ruhl [19] has also pointed out that the estimation of cooling rates

depended on the cooling conditions being Newtonian, intermediate, non-

Newtonian and ideal. He defined the condition of ideal cooling as the

state of ideal or perfect thermal contact between splat and substrate

giving rise to the condition of an infinite heat transfer coefficient

between the splat and substrate.

Under these conditions he reported that the cooling rate (T «

1/x ) where ~x' is the splat thickness. His calculations showed that

variations of splat temperature, solidification temperature and

substrate temperature had little effect on T compared with likely

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Table 1. Range of dendritic spacings associated with cooling rate and processes [7].

Cooling Ratt Characteristic Characteristic

Uniting Oendrltlc K/s Designation Processes Thickness* Spacing

£ 3 10 to 10 Very low Large sand cast- >6 m 0.S to 5 m

Ings and Ingots

10"* to 10° Low Standard castings 0.2 to 6 m SO to 500 vm and Ingots

lo" to 10^ Medium-High Thin strip, die 6 to 20 nm 5 to 50 ym casting, conven­tional atomlzatlon

103 to 106 High Fine powder atom- 0.2 to 6 nn 0.5 to S ym Izatlon, me It extrusion/ex­traction

10* to 103 Ultra-High Spray deposition, 6 to 200 va 0.0S to 0.S m and higher melt spinning,

electron/laser beam glazing

a. Assuno geometry is a metal slab from which the heat Is readily extracted from principal surfaces.

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Table 2. Summary of cooling rate measurements for splat cooling [12]

Location of measurement

Interface with substrate

Within splat

Exposed splat surface

Splat technique

Gun

Piston and anvil

Piston and anvil

Two-piston

Gun

{ {

Splat

Ag Al Au-14"/0Sb Fe alloy Al l'b Pb-50w/o Sn Al-4w/0 Cr

Al Cii and Cu-Ni

} } }

Substrate

Ni/Afj

Fused silica

Fe/Al

Mica on Cu

Cu

{

Thickness of splat s (urn)

~ 1

~27J 701 44t 44J 65t

~20

{

Initial cooling rate 'i'(Ks ')t

1 xl0"-5xl0« 1-5 x 10'-3 x 10' ~10s

up to 10" 6-9 x 10* 1-2 x 10* 7-8 x 10s

9 7 x 10'

^ up to 1-5 x 10*

} }

up to 10s

Reference

Prcdecki et al

Hrower ct al [18]

llarbur el al [14]

Miroshnichcnko and [15] Zakharov

Liililierj; and Miillcrf 15^

t Cooling rate '1' was measured tlicnnoclcctrically cxccpt by Hrower ct al and Lfihberg and Mi'tllcr who used optical pyrometry. llarbur el al and Lfihberg and Miiller were able to splat cool under an inert atmosphere.

} Half thickness (twin substrates).

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variations in ~ h1 and' s1 where 'h' is the heat transfer coefficient

between the splat and the substrate and ~s' is the splat thickness.

References[12,20-24] also provide the reader with an excellent

comparison of the cooling rates predicted by theory and the values

determined experimentally.

There are several techniques of RSP described in the

literature. Fig.6 shows a representative set of liquid quench [25]

procedures all closely related, which are commonly used in laboratory

preparation of metastable alloys. A description of the various

processes and their operation can be found in referencesf22,25-33].

Each of these processes result in the formation of flakes, foils or

tapes. Table 3 lists these processes and compares the cooling rates and

the final products obtained in each of the process. The methods

illustrated in Fig.6 are generally used to produce short experimental

specimens in either rod, strip or flake form. However, some of them,

for example the processes shown in Fig.6(d), 6(e) and 6(f) can however

be adapted to produce uniform continuous lengths of rapidly chilled

fibre or ribbon, or a continuous stream of staple fibre or powder like

material. Melt spinning and melt extraction, the relative quench rates

of which have been discussed by Pond and Winter [34] are two such

modifications. There are several other fiber processing techniques

illustrated in Fig.7. These methods are classified according to the

ways in which liquid metal is supplied to, and heat removed, from, the

solidification surface[35-39]. Table 4 lists these processes in the

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Figure 6. Schematic representation of various techniques of RSP.[25]

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Table 3. Quench rates and sample shapes in splat quenching [25].

Me thod

Cooling rate (#C sec"*) Sample shape

I Gun

Piston and anvil

Plasma-jet

(Torsion catapault

(Drum quenching

Roller quenching

106 to 108

106 to 107

~ 10

~ 106

~ 10€

~ 10

Irregular thin film or powder (rather uniform thickness)

Irregular thin plates (uniform thickness)

Irregular thin film (uniform thickness)

Irregular thin plates (uniform thickness)~

Regular uniform ribbon filaments

(uniform thickness)

Regular thin plates (uniform thickness)

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FREttUW

(a)

(c)

(e)

metal roo

QUENCHING MEDIUM AREA

souo wme

WATER-FILAMENT

(b)

(d)

( f )

Figure 7. Various techniques of filament casting: (a) melt drawing sheathed In glass; (b) melt extrusion Into a stabilizing medium; (c) melt spinning from a rotating chill surface;(d) melt-rolling (twin-roll quenching); (e) melt extraction from a crucible; and (f) melt-drag from a crucible side nozzle.[35]

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Table 4. Fiber Processing Techniques

Process Cooling Rate Productof solidification

1.Taylor's Wire 104-106K/s wires 2-50/im dia. Process[36]

2. Melt 105-108K/s filaments 1-100/im. Spinning [34] thick.

6 T 3.Rotating 10 -10 K/s ribbons 3-Bim. Cup[36]

4.TWin Roller[30] 105-10?K/s 10-100/m. strips. 5 8 5.Pendant 10 -10 K/s filaments. 25-125/jm. .

drop melt thick. extraction!36]

6.Melt drag[36,39] 102-104K/s. wire 0.1-8 mm. thick.

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order of cooling rates and the resulting product of solidification

obtained in each of the process. The methods described till now are

used for the synthesis of foils tapes or filaments. Similarly there are

various powder production processes two of which are illustrated in

Fig.8[40-43]. Table 5 in a similar fashion lists the processes and

compares the cooling rates obtained in each of the processes.

2.2 III-V Semiconductors

The III-V semiconductors are formed by combining atoms from

group IIIA and group VA. These crystalline semiconducting compounds

are not alloys but definite chemical compounds with a 1:1 atomic ratio

between the III and V atoms which occupy alternate sites in the crystal

lattice. Table 6 lists some of the important properties of the III-V

compounds[44]. Nearly all the III-V compounds crystallize

into two forms of crystal structures: "Zinc blende" which is cubic; and

"wurtzite" which is hexagonal.[45]

2.3 Ternary III-V Antimonides

Combinations of the six IIIA and VA elements give the

antimonide compounds listed in Table 7. Ternary solids of the

antimonides can be prepared by melting corresponding binaries of these

antimonides; since they are mutually soluble in all proportions with

the exception of GaPSb and InPSb. [46] For most of the ternary

antimonides, a solubility rule derived by Foster and Woods[l] is valid

which states that complete miscibility should be expected in all of the

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INDUCTION COIL

MELT

PRESSURE

l I I a • III i • j | I » • f't "1-J i •I'trr

CRUCIBLE

SPINNING CUP

TO DRIVE MECHANISM

1 k V V \ V V \ A

ORIFICE CHUOBUE

UOUIO / SOLID INTERFACE

MOLTEN DROPLETS

Figure 8. Powder preparation techniques: (a) Schematic arrangement of Rapid Spinning Cup (RSC); (b) Schematic arrangement of Twin Roller Atomization (TRA).[43]

M IO

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Table 5. Powder Processing Techniques

Process Cooling Rate Particle size

1.Centrifugal 104-10*K/S

Atcmization[ 41 ]

2.Spark Eroslan[42] 10#-10#K/s 5-15/on

3.Supersonic gas

Atcmlzatlon[43] >10*K/S <20/m

4.Rapid Spinning >10eK/s <50jum

Cup[43]

5.TWln Roll 10tt-10#K/s <30fzm

Atcnlzaticn[ 43 ]

6.Plasma Spray 107K/S -

Deposltlon[43]

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Table 6. Selected properties of III-V compounds, germanium and silicon [90].

Band Gap at Mdttef Point Vapor Pmui Material R.T. (ev) •o at M P. (aim) Cryul Strucnm

Si I.I 1 1420 10"* Cubic (diamond) Gc 0.67 936 I0"» Cubic (diamond) IN IO(Calc)»

— High Cubic (linc-bJcf«tc)

IN 4.6 (Calc)» >27001 High Ikugonal (wurohe) AIN >V >2900 >» Heaagtmal (wurtlite) CaN 3.2V >17001 >200* Hexagonal (wurtzite) InN U(Caic)> High Hcaagonal (wurtsite) BP i.V >ixxy >24» Cubic (linc-blcnde) AIP 2.42* >2000* High Cubic (mc-bknde) GaP 2.23 1467* »• Cubic (tinc-blcnde) InP IJJ 1060 23* Cubic (ainc-blcnde) BAa U(Cak)> >iooy High Cubic (ainc-blende) AiAi 116 >1700* High Cubic (tine-blende) CaAa 1.40 1231 09 Cubic (zinc-blende) InAl 0J3 Ml ass Cubic (nnc.btende) AlSb 1.74* ioso> <0.02 Cubic (awe-blende) CaSb 0.10* 706* <«x io-« Cubic (line blende) InSb 0.17 323 <IO*» Cubic (ainc-bhnde) InBi MctalKe 107 Vvrbo Ihnnona TIBi Metallic 230 Mnrhw Cubic oa

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Table 7. Possible combinations of the III-V Antlmonldes [1].

BINARIES! TERNARIES QUATERNARIES

AIIIBVA"W aiiibii:cv aiiIbIIiciii dv aiiibw ww A ° y 1-y x 1-* x y 1-x-y x y 1-x-y x 1-x y 1-y

AlFSb AlGaPSb*

AlSb AlAsSb AlGaSb AlGalnSb AlPAsSb* AlGaAsSb

udSb GaPSb InGaSb GaPAsSb AllnPSb*

InSb GaAsSb AllnSb InPAaSb AlZnASSb

InPSb GalnFSb

InAaSb GalnFSb

Systems which have not been prepared yet.

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systems where the lattice parameter difference is less than about 7.595.

However there are some exceptions as shown in Table 8.

2.4 III-V Ternary phase diagrams

Ternary III-V phase diagrams have three properties,

namely:

1) They are very useful in the LPE growth of III-V allays.

2) They are simple, and the liquid and solid solutions are

described well by simple solution models[47]. The simplicity of the

phase diagram is highlighted by the fact that the solid phase is

pseudobinary; that is, it has only one compositional degree of freedom,

and 3) The liquidus surface is simple. One liquidus surface, in

equilibrium with the pseudobinary solid, covers virtually the entire

ternary phase diagram.

The In-Ga-Sb ternary phase diagram has been determined

experimentally and has been verified by calculation. Further details

with regard to the calculation and evaluation of the phase diagram can

be obtained from references [48-59]. The usual ternary phase diagram is

a projection of the prismatic diagram upon a triangular base of the

prism. The liquidus isotherms are the intersection of the liquidus

surface and the isothermal planes. The solidus can be gained by a

similar way. The vertical plane intersecting at the two stoichiometric

binary compositions is called as "pseudo-binary" phase diagram. It

represents equilibrium between the stoichiometric liquid and solid. The

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Table 8. Lattice parameters of binary constituents in III-V ternary antimonldes [1].

Binary pair Lattice parameter difference (nn)

(*)

AlSb-GaSb 0,00 UO ;0,65

AlSb-InSb 0,03itU 5^5

GaSb-InSb 0.038U 6,11

InAs-InSb 0.0U21 6,72

GaAs-GaSb 0,0W»2 7,52

AlAa-AlSb 0.0U7U 8,02

IaP-InSb 0,0611 9,88

GaP-GaSb 0.06UU 11,15

AXP-AISb 0,0673 11,65

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"tie-lines" connect a point on the liquidus surface with the point on

the solidus surface with which it is in equilibrium. In the pseudo-

binary phase diagram the tie line connects the pseudo-binary liquidus

to the pseudo-binary solidus. In real cases the three III-V binary

compounds form an isomorphous system consisting of a liquid solution

(melt) and a solid sol\*tion. These two single phase regions are

separated by a three dimensional two phase region, which is bounded by

the liquidus surface at the top and the solidus surface at the bottom.

The three binary phase diagrams are the side - faces of the prism. Most

of the III-V ternary compounds have liquidus and solidus surfaces

characteristic of ideal systems.

The InSb-GaSb pseudo-binary phase diagram was determined by

Blom and Plaskett [49] as a special case from the ternary system where

the antimony fraction in the liquid remains constant at 0.5 and

therefore in the solid also. However for these calculations they found

that the pseudo-binary liquid is independent of the interaction between

In and Sb and Ga and Sb. Therefore, in general, the pseudo-binary

liquid can in first order be described by the interaction between the

two metals and is independent of the group V element. The pseudo-binary

phase diagram for InSb-GaSb system is shown in Fig.9.

2.5 Synthesis Of Bulk Crystals Of Ir fiaj Sb:

Primary investigations in this system were reported by Russian

workers Koster and Thoma [60] and Goryunova et al. [61] wherein they

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900

InSb GoSb

Figure 9. The InSb-GaSb pseudoblnary phase diagram showing O :experimental data;*: differential thermal analysis; X ordinary thermal analysis; A : x-ray measurements; Q : [49]

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tried to obtain single phase alloys as a result of progressive

annealing of powdered alloys at high temperatures. They tried to

achieve this with several compositions in this system. In later work

Woolley and Lees [62,81] showed that complete solid solution can be

obtained in the GaSb-InSb system provided that the alloys are annealed

for a sufficiently long period of time at the correct temperature. The

procedure followed was sealing the required quantities of the

appropriate compounds in evacuated sealed quartz tubes, heating to a

temperature above the melting point of the higher melting point

compound, and quenching in water. The resultant alloy was then powdered

and X-rayed following which the alloy would be again compressed and

resealed in vacuo and annealed to a temperature corresponding to the

melting point of the lower melting point component initially and then

progressively annealing it at higher temperatures depending on the X-

ray analysis until equilibrium was reached. However exceedingly slow

diffusion rates in the materials required long annealing periods

extending upto several weeks and in extreme cases several months before

the system showed any movement towards an equilibrium condition. The

problem was more severe in the case of solids as compared to the

powders of the same composition. The sluggishness of the system is

attributed to the very slow diffusion rates of the individual elements

in the solid state. Boltaks [63] studied the diffusion of these

elements in detail.

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Ivanov-Omskii and Kolomiets [64] on realizing the difficulties

faced in the preparation of equilibrium solid solutions of these

compounds, used zone melting to achieve equilibrium conditions. Zone

levelling was also reported by McGroddy et al.[65] and Blakemore[66].

McGroddy et al.[65] reported that the ingots were polycrystalline with

grain size too small to allow preparation of single crystal samples.

Plaskett and Woods[67] reported the preparation of

polycrystalline alloy ingots - covering the entire composition range by

zone levelling techniques. They used zone travel rates of the order of

1 mm/hour. They examined the alloy metallographically, and obtained

lattice parameter values using X-ray diffraction. They also used

optical transmission techniques to determine the band edge. Joullie et

al.[56] reported the preparation of n-type single crystals of ternary

In aj.jjSb with 0<x<0.45 by means of a solution growth method (TGZM)

using indium as solvent and indium antimonide as seed. The charge

sandwich constituted polycrystalline Ga n j b obtained from a

Bridgman ingot. TGZM growing method consists in setting a seed charge

sandwich in a temperature gradient of the order of 80-

140°C/cm[68,69,70]. The charge which is dissolved, diffuses across the

liquid and crystallizes on the seed, while the solvent zone travels

towards highest temperat\ire regions. Growth and characterization of

solid solutions by TGZM was also performed by Hamaker et al.[68].

They obtained single crystal layers of upto 4 mm in thickness with

cross sections of 0.7 x 0.7 cm using <111> seed orientations of InSb at

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a growth temperature of 420°C. Similar studies using TGZM was also

carried out by Michel et al.[71].

Solute diffusion growth for the preparation of ternary In Ga- ..

xSb crystals was reported by Miller et al.[72]. In this method a metal

alloy of a selected composition along with a quantity of the metalloid

element is placed in an evacuated sealed silica growth ampoule. The

growth ampoule is then placed in a growth furnace when the metalloid

will slowly distill or sublime and will enter the metal solution by

vapor transport at a rate that can be controlled by controlling the

metalloid reservoir temperature, Tc and its relationship to the

solution temperatures. Rodot et al.[73] also applied this method for

preparing sound ingots of Ga In .j Sb.

The travelling heater method (THM) of crystal growth, whereby a

solvent zone is caused to move through a solid feed material by the

movement of a heater; and first conceived by Pfann[69] in 1955 was

applied by Yip et al.[74,75]. They prepared cylindrical feed rods of

GaxIn2_xSb by gradient freeze technique in a vertical furnace[79].

Directional solidification of InxGa1_xSb was studied by Wilcox et

al.[76,77], They prepared ingots on the earth and in skylab, and

compared the influence of gravity on the crystal defect formation.

Joullie et al.[78] have also reported use of the vertical freezing

technique for preparing bulk GaxIn1_xSb using travel rates of the order

of 0.25-1 mm/hour. J. F. Yee [79] and R. A. Lefever[80]have also

reported preparation of bulk crystals of GaxIn1_xSb using the

travelling heater method.

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Bachmann et al.[82] have succeeded in growing radially

homogeneous bulk single crystals using the czochralski method. They

synthesized single crystals of GajjIn -j Sb in the range 0.9<x<l.

2.6 Thin Film Methods

Crystalline and amorphous films of In Gaj Sb have been made by

several techniques such as liquid phase epitaxy, gas phase epitaxy,

vapor phase epitaxy, molecular beam epitaxy, flash evaporation and

multitarget sputtering. Greene et al.[83] have reported the growth of

single crystal epitaxial InxGa1_xSb thin films at compositions ranging

from x = 0 to x = 1, on a variety of substrates (BaF2, NaCl and InSb)

using multitarget sputtering.

Yano et al.[84] have reported growth of In Gaj.j Sb by

molecular beam epitaxy. They identified the relation of the arrival

rates of In and Ga molecules to the composition of the epilayer. The

films were grown on (100) surface of Cr doped semi-insulating GaAs.

Gotoh et al.[85] studied the growth of In0#2Gao.8sb on Cr doped

(100) GaAs using H2S gas for sulfur doping to change the conductivity

type.

Single phase InxGa1_xSb films by flash evaporation have been

prepared by Syamalamba et al.[86,87]. They prepared the films of

several compositions by controlling the deposition parameters and

characterized them for structural, compositional and optical

properties. Amemiya et al.[88] reported electrical property

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measurements on films of ternary In Gaj Sb grown by a source

temperature programmed vacuum evaporation technique which is a

modification of a simple two temperature evaporation technique.

Wickersham et al.[89] carried out impulse stimulated "explosive"

crystallization of sputter deposited amorphous films of (In,Ga)Sb.

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CHAPTER 3

EXPERIMENTAL SET UP AND PROCEDURE

3.1 MATERIALS PREPARATION

Preliminary experiments

Preliminary experimental work was performed by weighing

stoichiometric amount of InSb and GaSb required for obtaining the

desired composition in quartz tubes. The tubes were of standard wall

thickness and 0.3 cm. inner diameter. The tubes were batched with

99.999% InSb and 99.99% GaSb obtained from a commercial source1. The

samples were obtained in the form of 0.3 cm. pieces which were ground

to a powder in order to seal them in the tubes and batched in a glove

box back filled with nitrogen gas to prevent any atmospheric

contamination. The quartz tube was batched with InSb at the bottom

followed by GaSb, as shown schematically in Fig.10. The tubes were then

sealed after evacuating in a vacuum of 10~5 torr and heated in a

resistance heated vertical tube furnace to a temperature of 900°C.

The furnace temperature was measured with a thermocouple set

inside the furnace. The tubes were then held at that temperature for

1. Cerac Chemicals, Box 1178, Milwaukee, Wisconsin 53201.

35

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Fused Silica Tube

Hot Zone

A

^

tO"5 Torr

In £b

— Go. Sb

Figure 10. Schematic diagram of synthesis of In Ga, Alloys. 1

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eight hours to ensure good homogeneity and then quenched into a bath

comprising of a mixture of ethylene glycol and dry ice giving a final

temperature of -31°C. The quenched ingot gave sharp crystalline X-ray

diffraction peaks and hence it was decided to pursue other rapid

solidification techniques to obtain higher quenching conditions.

Initial test run-was conducted on a melt spinning set up which

.consisted of melting the alloy in a graphite susceptor placed in a

quartz tube. Ejection of the melt onto the wheel resulted in the

production of flakes. In order to collect these flakes, it was decided

to adopt the spinning cup mode of rapid solidification.

Design Of The Spinning Cup

The spinning cup basically comprised of two parts. An outer cup

made of aluminum, 17 cm. in diameter, and a disk made of copper, 0.3

cm. in thickness and 15.3 cm. in diameter. The copper disk was placed

at the bottom on the inside of the aluminum cup. The entire set up was

6.5 cm. in height and the aluminum cup was provided with a 5 cm. hub

which was fixed onto the shaft of a AC/DC motor capable of 10000 rpm at

0.5 HP with the help of set screws. In this way the spinning cup was

fixed to the motor. The motor speed was varied with the help of a

variac. This set up along with the motor was then placed inside a box

made of plexiglass covered on all four sides except at the top. This

arrangement was then placed close to the resistance heated vertical

tube furnace used earlier for tube quenching. Fig. 11 shows a schematic

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Argon flow

Quartz tube ~ 10x12 mm

Copper disk Vfe in.

.V2 HP motor 10,000 RPM max.

Figure 11. Shanatlc of the 1RSP set up.

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of the set up described above.

Melting and Ejection

The melting and ejection step involved melting the prealloyed

ingots of the desired composition. These ingots were made as described

earlier in the tube quenching experiments except that instead of

quenching; the tubes were allowed to cool in the tube furnace. The

tubes used in this spinning cup arrangement were 17 cm. long and were

provided with a nozzle. Two nozzle specifications were tried out in the

experiments, one with a diameter of 0.1 cm. and the other with a

diameter of 0.15 cm. Hie top end of the quartz tube was connected to a

solenoid valve by copper and tygon tubing. The solenoid valve was then

connected to the argon gas tank. The desired pressure of the argon gas

was obtained by adjusting the pressure in the regulator.

The melting sequence consisted of heating the furnace to the

desired temperature of 700°C or 900°C depending on the superheat

requirements. The quartz tube with the sample was then inserted into

the furnace to a depth of 10 cm. from the top of the furnace. Prior to

the introduction of the tube into the furnace a high flow of argon was

passed through the tube to remove the trapped air inside. The tube was

held in the furnace at the required temperature for a period of 30

minutes. Meanwhile the cup was fixed onto the motor and allowed to spin

at the desired rpm. The tube was then removed and the liquid metal

immediately ejected onto the rapidly rotating disk inside the cup, with

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the help of the solenoid valve.

A large number of trials and experiments were carried out with

several compositions of the alloys at various rpm settings of the motor

using tubes provided with nozzle diameters of 0.1 and 0.15 cm. The rpm

of the cup was calibrated to the variac setting using a stroboscopic

light. In this way the cup was made to rotate between 4000-10000 rpm.

The flakes and powders were analyzed using X-ray diffraction, optical

microscopy, SEM, TEM and Auger electron spectroscopy for characterizing

the microstructure. Phase transformation investigations were conducted

using the DSC.

The flakes and powders were examined as is for topographical analysis

and etched to reveal the microstructures. The surface as well as the

cross section of the powders and flakes were examined.

Heat Treatment

The as-prepared flakes were sealed under a vacuum in quartz

tubes and heat treated in a drive in furnace to temperatures of 100,

200, 300 and 400°C for a period of two hours at each of the

temperatures to induce annealing which are discussed in the later

chapter. The samples were then again mounted in casting resin and

polished on which the electrical resistivity measurements were made.

3.2 MATERIALS CHARACTERIZATION

X-ray Diffraction

X-ray diffraction analysis was performed on the as-prepared

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flakes and powders on a standard diffractometer using monochromatic

copper Ka radiation provided with a nickel filter. The beam slit used

was 1°medium resolution (MR) and the detector slit used was 0.2°MR. The

furnace cooled ingots were also subjected to X-ray analysis by crushing

to -325 mesh size fractions. The instrument was calibrated for the

proper 28 values using an internal standard of high purity silicon.

For all the runs a scanning rate of two degrees per minute was used.

Optical Microscopy

Optical microscopy was done on the flakes and powders which

were etched to see the grains and the columnar structures during the

solidification of the alloys. This was done by mounting the flakes in a

casting resin to see the morphology on the surface. In order to see the

cross section, the flakes and powders were sandwiched between the

casting resin and then cut at right angles to see the cross section of

the solidified flakes and powders.

Scanning and transmission electron microscopy

Scanning electron microscopy (SEM) and energy dispersive X-ray

analyses (EDAX) analyses were performed on the as-prepared powders and

flakes. The flakes and powders were mounted on standard sample stubs

and sputtered with Au/Pd. The etched samples were mounted in a casting

resin, and the entire resin surface excluding the sample was coated

with colloidal carbon to prevent charging of the electron beam.

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Transmission electron microscopy(TEM) was performed on crushed

powdered samples dispersed in n-batanol and floated on carbon coated

copper grids. Selected area diffraction was used to determine the

amorphous or macrocrystalline nature of the alloys.

Scanning Auger Electron Microscopy

Scanning Auger electron microscopy was done on the as-prepared

flakes to see if there was any oxidation and to see if there was any

formation of oxides. This was done by taking area scans and depth

profiling by argon sputtering at 2keV using an electron beam of 3keV.

Thermal Analysis

Differential scanning calorimetry (DSC) was performed on as-

prepared flakes and powders. This was done to see the effect of rapid

solidification on the structure and to perceive any phase

transformations if any. The analysis was performed in a Dupont 1090

Thermal Analysis System,2 which is computer interfaced allowing fast

real time data acquisition at maximum sensitivity.

The software automatically determines peak positions. The as-

prepared powders and flakes (typically weighing 10-20 mg) were sealed

in copper pans; a small opening was provided in the pans to prevent any

pressure build up in case of volatile elements being present in the

samples.

2. Dupont 1090 Thermal Analyzer, E.I. Dupont and deNemors Inc., Wilmington, Delaware.

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The instrument could record any thermal changes up to a temperature of

600°C. A heating rate of 10°C/min was used for all the samples, and a

neutral atmosphere of nitrogen gas was maintained in the DSC cell.

Electrical Resistivity Measurements

A four point probe was used to measure the resistivity of the

heat treated flakes. The samples were annealed for two hours at

temperatures of 400,500,600 and 700°C, respectively, and then mounted

in casting resin and polished. Hie measurements were performed at room

temperature after making the necessary correction for the sample shape

and thickness. These corrections were obtained from reference[97].

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CHAPTER 4

RESULTS AND DISCUSSION

4.1 X-ray Diffraction Analysis

The X-ray diffraction analyses were performed on all rapidly

solidified samples which are discussed separately in two subsections

entitled General Analysis and Specific Analysis.

General Analysis: This section includes the analysis performed on the

tube quenched ingots and the analysis on the resultant flakes and

powders obtained by melt spinning initial trial melts. The compositions

investigated and the melt spinning conditions are listed in Table 9.

Tube quenching experiments were performed on alloys of composition

In0.53Ga0.47Sb' In0.44Ga0.56Sb' In0.46Ga0.54Sb' and In0.41Ga0.59Sb"

X-ray diffraction analysis performed on all these samples

indicated that the samples were strongly crystalline, and it was thus

necessary to pursue the alloy processing by melt spinning to induce

microcrystallinity or amorphism. All the compositions listed in Table 9

were initially ejected onto the Cu disk inside the spinning cup(see

Fig.11) at a distance of 2.54 cm. from the center. At a relatively low

speed of 4000 rpm, flakes were obtained and at 6000 rpm, a combination

of flakes and powders were obtained as shown in Table 9. The X-ray

traces of the ingots for each of the above compositions were used to

44

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Table 9. Melt Spinning Compositions with Melt Spinning Parameters

Composition T melt rpm Velocity Pressure Nozzle Ejection run Form Remarks

(°C) (m/s) x 104Pa diameter distance (cm) (cm)

.ls .e?813 960 6900+100 18.4 7 0.09 2.54 1 Flake -

In0.13Ga0.87Sb 967 10000 +100

27.26 7 0.09 2.54 2 Flake -

In0.16Ga0.84Sb 940 8000+100 21.3 8 0.09 2.54 3 Flake -

966 5600+100 14.62 7 0.09 2.54 4 Flake -

860 8000+100 21.3 2 0.09 2.54 5 Flake powder

-

700 9000+100 23.93 4 0.09 2.54 6 Flake powder

-

In0.51Qa0.49Slb 865 6500»+100 17.3 4 0.09 2.54 7 Flake -

835 9500+100 25.27 2 0.09 2.54 8 Flake powder

-

o.si o 813 820 9500+100 25.27 4 0.09 2.54 9 Flake powder

melt held for 3 sec.

Continued

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Table 9. Melt Spinning Compositions with Melt Spinning Parameters

Composition T melt rpn Velocity Pressure Nozzle Ejection run Form Remarks

(°C) (m/s) x 104Pa diameter distance (an) (cm)

825 9500+100 25.27 4 0.09 2.54 10 Flake powder

melt held for 18 sec.

In0.54Ga0.46Sb 700 9500+100 25.27 2 0.09 2.54 11 Flake powder

cup fixed vertical

.ga o.o?813 700 10000 +100

26.6 4 0.09 2.54 12 Flake powder

-

710 9500+100 25.27 4 0.09 2.54 13 Flake powder

-

In0.446a0.56Sb 680 10000 +100

26.6 4 0.09 2.54 14 Flake powder

-

In0.44Ga0.56Sb 720 9500+100 25.27 2 0.09 2.54 15 Flake powder

-

In0.44Ga0.56Sb 720 8000+100 21.3 1 0.09 2.54 16 Flake powder

-

In0.44Ga0.56Sb 830 10000 +100

26.6 2 0.09 2.54 17 Flake powder

-

Continued

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Table 9. Melt Spinning Compositions with Melt Spinning Parameters

Composition T melt rpn Velocity Pressure Nozzle Ejection run Form Remarks

(°C) (m/s) x 104Pa diameter distance (cm) (cm)

In0.44Ga0.56Sb 749 4250+100 33.91 4 0.15 2.54 18 Flake 3 sizes length : > 5mn.l8A = 5mm.18B < 5mm.18C

In0.44Ga0.56Sb 735 6500+100 51.86 4 0.15 2.54 19 Flake 19F powder 19P

In0.44Ga0.56Sb 735 8020+100 63.99 4 0.15 2.54 20 Flake 20F powder 20P

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perform lattice parameter calculations, summarized in Table 10.

These calculations provide evidence that the alloys do not

yield an equilibrium solid solution upon solidification. In addition,

some ingot compositions show presence of two non- equilibrium phases

and on subsequent melt spinning there is either a retention of one of

these phases or a different non-equilibrium phase formation as

indicated in Table 10. In the others the non-equilibrium phase is

retained on quenching. A common feature however is a noticeable

disappearance of peaks in the traces obtained for the flakes and

powders indicating disordering as shown in Fig.12 which is a

representative trace for alloys of composition Ino.44Gao.56s 3, A so the

X-ray traces for the powders are relatively broad in comparison with

those obtained for the flakes; this is a result of smaller crystallite

sizes obtained in the case of the powders, because cooling rates in the

powders were probably greater than the cooling rates in the flakes.

These data suggest that the powder is probably formed in the present

processing by a mechanism similar to that described by Glickstein[41]

wherein the liquid metal on contact with the disk disintegrates into

molten droplets which are solidified either in air or on impinging the

walls of the cup. This is the mechanism for powder formation in the

centrifugal atomization technique described in reference[ 41 ].

Because of the large surface area associated with these

particles there is rapid cooling resulting in high cooling rates. The

broadening of the peaks which is induced due to reduction in

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Table 10. Lattice Parameter Values

Composition Lattice Parameter A°

Literature value [44]

Ingot Powder Flake Structure Literature value [44] a2

.IS .S?813 6.15 6.095 - - 6.095 diamond

lh0.1GPa0.84sb 6.16 6.095 - 6.095 diamond

In0.43Ga0.57Sb 6.265 6.447 6.29 6.23 6.23 diamond

In0.44Ga0.54Sb 6.27 6.095 6.29

'6.095

6.29

6.095

diamond

6.275 6.447 6.29 6.192 - diamond

In0.51 0.49s*5 6.295 6.447 6.29 6.29

6.095

6.29 diamond

In0. . 07Sb 6.48 6.447 6.447 6.447 diamond

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Powder

(III) (220)

(311) a •f/AN—, ,J T-+-I : 1—I—I

Flake (/)

w-O

(220)

C

Ingot

(220)

(311) (422) uu (331)

50 48 ANGLE (20 degrees)

Figure 12. X-ray trace of ingot, flake and powder obtained from melt of run 19.

8

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crystallite size was observed only in powders obtained from melts of

initial composition InQ 51GaQ 49Sb and In0 44Ga0 56Sb. This suggests

the marked influence of RSP only on these two compositions. X-ray

traces for powders obtained for other compositions showed sharp

crystalline peaks further supporting the above inference.

Specific Analysis

Specific X-ray diffraction analyses were conducted on flakes

arc! powders obtained for alloys of initial composition Ino.5iGao.49Sb

and In0-44Ga0 56Sb.

Based on previous work of others [ 2 ], experiments were designed

to assess the effect of undercooling on the structure by reducing the

superheat. The experiments were carried out by removing the quartz

tube from the furnace after the alloy was melted and homogenized for a

half hour at 800°C (a superheat of 160-200°C) and holding the tube for

several seconds in order to decrease the temperature before ejecting

the alloy. Three experiments based on this procedure were conducted on

the Ino.5iGao.49st> composition. The experimental parameters are listed

in Table 9. They have been identified by runs 7-10. All these runs

resulted in the production of flakes and powders which exhibited

similar X-ray traces. Lattice parameter values for these alloys are

shown in Table 10. Again it is clear that there is formation of mixed

phases and a retention of one of the phases in the flakes and powders

along with a marked broadening effect of the peaks for the powders.

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In all the above experiments the parameters varied In melt

spinning were the holding times, temperature and the disk velocity.

From the X-ray analysis it was concluded that the velocity of the disk

and the redaction in superheat played a significant role in optimizing

the effects of rapid quenching. Therefore melt spinning of

In0.44Ga0.56sb was performed at a reduced but constant superheat of

100°C while only varying the disk velocity. Further, all the

experiments on this alloy were performed by ejecting the liquid metal

at the periphery of the copper disk to obtain maximum peripheral

velocity to act on the metal so that along with the centrifugal force

effect there would also be shearing of the metal droplets inducing more

effective quenching conditions.

These melts are identified by runs 18, 19, and 20. Run 18

resulted in big flakes ( > 5 mm length), small flakes (=< 5mm length)

and fines (< 5mm length ) identified as 18A,18B and 18C, respectively,

in the discussion below, whereas runs 19 and 20 resulted in the

production of flakes (identified as 19F,20F) and powders (identified as

19P,20P). The X-ray traces for all these melts showed essentially

similar peaks to those obtained for Ir\).44Qa0.56s'3 sh°wn in Fig. 12 for

run 13. Runs 13 and 19 showed identical X-ray traces. The X-ray traces

however showed a relatively larger amount of peak broadening in

comparison to the previously mentioned compositions. This was because

the experiments were performed by ejecting the liquid metal on the

periphery of the disc.

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In order to see the effect of heat treatment on the rapidly

solidified flakes, X-ray analyses were performed on the flakes of

composition In0.5iGa0<49Sb obtained from run 7, after they were

subjected to an annealing treatment at temperatures of 100, 200, 300

and 400°C for a constant period of two hours. The X-ray trace obtained

for the flakes annealed at 400°C was identical to the trace obtained

for the ingot. In the cases of the samples annealed at temperatures

0 other than 400 C, the traces indicated a gradual increase in the

intensities and the reappearance of peaks, which were absent in the

quenched flakes.

In summary, it is reasonable to conclude that the flakes and

powders contain a mixture of metastable phases, similar to the

unquenched ingot, and show a general broadening of peaks, which

suggests very fine grains due to rapid solidification. Flakes obtained

from run 7 retained one of the phases seen in the unquenched ingot

(Table 10); however heat treatment of these flakes results in the

reappearance of the original phases present in the initial ingot.

4.2 Surface Morphology

The as-prepared flakes and powders were examined under the SEM

using EDAX to observe the morphology of the flakes. Fig. 13 shows

micrographs of surfaces of the flakes away from the copper substrate.

It appears that liquid metal droplets have solidified on the surface

to give it a very irregular morphology. At high magnification these

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Figure 13(a). SEM micrograph showing surface of flake obtained from run 9, mag. 33 X.

Figure 13(b). SEM micrograph showing surface of flake obtained from run 20, mag. 1500 X,

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Figure 13(c). SEM micrograph showing precipitates in flakes obtained from run 1, mag, 4200 X.

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solidified droplets appear like precipitates (see Fig.l3(c)).

SEM(EDAX) was performed on the precipitates and on the matrix

in order to study any variations in composition. The precipitates

showed a considerable decrease in the Ga counts in comparison to the

matrix.(see Fig. 14 and 15) Hence it can be concluded that the

precipitates are richer in InSb as compared to GaSb. These precipitates

were observed in all the compositions. However these precipitates were

only seen on the surface of the flakes solidified away from the copper

substrate. The surface of the flake adhering to the copper disk showed

a much smoother morphology in comparison to that discussed above. There

was also no compositional variation.

The powders were also examined voider the SEM. Fig.16 shows the

morphology of the powders in a series of micrographs taken on powders

obtained from run 8, 9, 10 shown in Table 9. It can be seen that the

powders are irregular with particle sizes ranging from 1 micron to

several 100 microns. A noticeable feature is the cluster of spherical

particles in photomicrograph 16(d) with particle sizes being 2-3

microns. Micrograph 16(e) shows the fractured part of a large powder

particle. The fractured surface shows the columnar growth indicating

the directional mode of solidification. Micrographs 16(f) and 16(g) on

the other hand do not show the formation of any spherical particles.

These were the powders obtained from run 8 and 9. The particles are

very irregular, and the sharp edges, in fact, indicate that they are

probably formed by large flakes fragmenting into smaller ones. However

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2888

COUNTS

PRECIPITATE ANALYSIS

ENERCT (KEVJ

Figure 14. EDAX analysis of the precipitate observed on the flake surface.

2

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IBM

MATRIX ANALYSIS

COUNTS

t.a te.e ENCRCY IKEVI

Figure 15. EDAX analysis of the matrix on the flake surface.

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Figure 16(a). SEM micrograph showing the powder obtained from run 10, mag. 360 X.

Figure 16(b). SEM micrograph showing the powder obtained from run 10, mag. 3600 X.

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SEM micrograph showing the powder obtained from run 10, mag. 1090 X.

SEM micrograph of powder obtained from run 10, mag. 1620 X.

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Figure 16(e). SEM micrograph showing fractured powder particle obtained from run 10, mag, 1620 X,

Figure 16(f). SEM micrograph showing the powder obtained from run 9, mag. 660 X.

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Figure 16(g). SEM micrograph showing the powder obtained from run 8, mag. 430 X.

*

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the particles occur in various shapes so that they could be occuring by

a combination of 'centrifugal atomisation' and by fragmentation.

4.3 Thickness Measurements

Figure 17 shows micrographs of cross sections of a flake.

Micrograph 17(a) shows a part of the flake surface in contact with the

copper disk, while 17(b) shows the surface not exposed to the disk and

hence exhibits a very rough surface in contrast to the relatively

smooth surface shown in micrograph 17(a). However both micrographs

indicate the columnar mode of solidification which is expected since

heat flow is basically directed towards the copper disk. The striations

on the surface seen in Fig. 17(a) are due to the flake being in contact

with the disk.

The thickness of the flakes were measured from micrographs of

one or two random flakes flakes obtained from melt spinning each of the

various compositions listed in Table 9. Thickness measurements were

made by taking into account the distortion in the secondary electron

image due to the tilting of the specimen. The values, estimated after

incorporating the correction were then correlated with the tangential

speed of the disk at the point where the liquid metal was ejected. A

plot of the flake thickness against the speed is shown in Fig. 18. Each

point on the plot represents an average value of 8-10 thickness

measurements. The bar around each point is a measure of the standard

deviation. The plot shows that the flake thickness decreases with

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Figure 17(a); SEM micrograph of cross section of flake obtained from run 1, mag. 1700 X,

Figure 17(b). SEM micrograph of cross section of flake obtained from run 19, mag. 770 X.

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Ti 5~TT VELOCITY fm/s)

¥ lo So" lo 80

Figure 18. Variation of flake thickness with disk velocity.

o> CJl

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66

increasing speed, but for speeds greater than 35m.s-1 a slight increase

in overall flake thickness is observed. This may suggest that some

flakes lose contact with the disk resulting in higher thickness

values. It may be noted that the plot may not be a true representation

of the entire process since the measurements were made on one or two

random flakes.

The thickness measurements were also used to estimate the

cooling rate after making a reasonable assumption that the heat

transfer coefficient at the interface between the flake and the

substrate is infinity, rendering conditions of ideal cooling to be

applicable. This assumption is applied from the fact that the majority

of the flakes exhibited excellent wettability!19]. Based on this

assumption, thecooling rate,

(CR) «=ci/d2 (2)

where CR = cooling rate from the initial splat temperature to the half

temperature, given by,

*1/2 = 1/2(TS+TB) (3)

where Ts = initial splat temperature and TB = initial substrate

temperature and d = thickness of the flake. Ruhl[19] made calculations

based on a one dimensional heat flow analysis and obtained a

proportionality constant of approximately 81 cm2°C/S which agreed quite

well with the analytical solutions obtained for the cooling rate from

Cars law and Jaegar[98]. Ruhl's analysis gave the following relation for

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the cooling rate,

CR = 81/d2 (4)

This square relation is reported to hold true for ideal cooling

conditions and is found to be marginally affected by changes in

solidification temperature, initial substrate temperature, and splat

and substrate material. Assuming the validity of the above equation, a

range of cooling ra.tes were estimated for the flake thickness

experimentally determined. Fig. 19 shows the variation of the estimated

cooling rates with experimentally determined flake thickness. As can

be seen cooling rates varied from 3.1 x 106 to 3.2 x 10 °C/s,

the higher cooling rates being obtained for flakes of diminishing

thicknesses. It may also be noted that thickness is also involved in

the parameter hd/kg (d=flake thickness, h=heat transfer coefficient and

kg=thermal conductivity of the flake) which determines whether ideal

cooling exists. As *d' becomes smaller, often the cooling will cease to

be ideal unless ~h' is very large. In our discussion since the flakes

made very good contact with the substrate we assume ideal cooling

conditions whereby, then the above square law holds. [19]

4.4 M icrostructure Analysis

The microstructure of the flakes and the powder were recorded

by optical microscopy and scanning electron microscopy after etching

them with an etchant consisting of the composition 2HF:2HAc:3HN03:6H20.

Initial composition of the etchant was obtained from the

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To Jo THICKNESS (MICRONS)

¥~ Jo 80

Figure19. Plot of experimentally determined flake thickness (d, microns) versus cooling rate (°K/s) calculated assuming the validity of the equation Cd =81 [19]

o» CD

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literature[76], however the etchant was very strong and had to be

diluted with water until finally, etching the flakes with the above

composition for 1 second worked extremely well in revealing the grain

structure of the rapidly solidified flakes and powders. The etchant

however failed to reveal any structure on the furnace cooled ingots.

The ingots on the other hand being bireflectant [ 79 ] show the grains by

reflection of light incident on a flat cross section of the ingot at an

angle. Fig.20(a) shows the micrograph of the cross section of the,

ingot with the grains being revealed by the property of bireflectance

exhibited by this alloy.

The figure shows random orientation of the grains since the

ingot was just allowed to cool in the furnace. Moreover the grains are

not uniform, being elongated more in one direction than the other. An

average grain size of 970 microns was estimated for the ingot by taking

an average of 13 measurements for several grains seen in Fig. 20(a).

Fig. 20(b) shows optical micrograph of the surface of the flake etched

with the above etchant. The flake was obtained by rapidly solidifying

the ingot of initial composition Ino.51GaQ.4gSb,represented by run 7 in

Table 9, the surface of which clearly shows the cellular grain

morphology.

SEM micrograph of the same area is shown in Fig. 20(c). The

pores seen in the micrograph correspond to the etch pits. Since the

topographical analysis of the flakes revealed presence of second phase

precipitates on the surface the etched surfaces were examined for any

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Figure 20(a). Optical micrograph showing the cross section of an ingot of composition Ino.46Gao,54Sb» ma9- 10

Fiqure 20(b). Optical micrograph showing the etched surface of the flake obtained from run 7, mag. 1000 X.

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SEM micrograph showing the etched surface of the flake observed in Fig. 20(b), mag. 2300 X,

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compositional inhomogeneities. A number of spectra were taken on the

grains and on the grain boundary to observe any changes. However no

differences in the counts for any of the components was seen; this

means that the precipitates formed only on the surface and subsequent

polishing during sample preparation eliminated them.

The cross section of the flakes and powders were then examined

to observe the structural morphology. Micrographs in Fig. 21 show the

etched cross section of the flakes. Fig.21 (a) shows the optical

micrograph of the flake of composition Ing gjGaQ 49Sb, obtained from

run 7 in Table 9, whereas Figs. 21(b),21(c) show the microstructure of

etched cross section of flakes of composition Ino.44Ga0.56sb obtained

from run 14 in Table 9. The micrographs clearly show the columnar

grains traversing from the top to the bottom which is the direction of

heat flow towards the copper disk.

Micrographs 21 (a),21(b), show the solidification morphology and

in some cases clusters of very fine dendrites. Fig.21(c) represents an

SEM micrograph of a section of Fig.21(b). The micrograph shows the

branching of the dendrites. The photomicrograph shows an irregular

morphology of the dendrites which may be because of the rapid rate of

heat extraction induced by the process. Another reason is the rotation

of the disk which causes shearing of the liquid metal.

There was no compositional variation seen in the cross section

of the flakes. The cross section of the powders also show similar

behaviour as the flakes. Figs. 22(a) and 22(b) show micrographs

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Figure 21(a). Optical micrograph showing the cross section of flake obtained from run 7 etched with the modified reagent, mag. 1000 X.

: * V* ' •

w <4 V-cm; f s -"•? * —

BR 3

• ,*w«. .. - •

' • J

-A v?**-« %

Figure 21(b). Optical micrograph showing the cross section of flake obtained from run 14, mag. 1000 X.

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SEM micrograph showing the cross section of flake observed in Fig. 21(b), mag. 1700 X.

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Figure 22(a), SEM micrograph showing the etched cross section of powder obtained from run 14, mag. 3700 X.

Figure 22(b). Optical micrograph showing the etched cross section of powder obtained from run 14, mag. 1000 X.

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t

Figure 22(c). Optical micrograph showing the etched cross section of powder obtained from run 10, maq. 1000 X.

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obtained from the cross section of powders of composition

InQ 44Ga0 5gSb obtained from run 14 in Table 9. The micrographs show

the columnar grains similar to those shown for the flakes. Fig.22(c)

shows the micrograph which represents the cross section of a powder of

composition Ino.5iGao.49sb represented by run 10 in Table 9. The

micrograph also shows columnar grains except that certain parts of the

powder remain unetched suggesting the presence of amorphous regions.

Similar evidence of unetched amorphous regions was also observed in

splat quenched Fe80B2oflakes investigated by Turnbull[25].

In order to see the effect of the cooling rate and disk vlocity

on the microstructure, flakes obtained from run 7, 18, 19, and 20 were

etched and examined under the optical microscope and the SEM. Figures

23 and 24 show the optical and SEM micrographs respectively for the

flakes obtained from runs 18,19 and 20. Grain size estimates were made

from each of the micrographs using the straight line method described

by Underwood et al.[91]. A plot of the grain size against the disk

velocity is shown in Fig.25. Hie plot exhibits a behaviour similar

to the trend seen in the thickness - velocity plot in Fig. 18. The

grain sizes decrease with increasing disk velocities but increase at

very high velocities due to insufficient disk-flake contact at such

high velocities. Therefore, an important conclusion from Fig. 18 is

that the data reflect a velocity band in the range 45-60 m/s where the

effect of rapid quenching is at its peak. This also suggests that glass

formation tendency is the highest in this velocity band for the

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Figure 23(a). Optical micrograph showing the etched surface of the flake obtained from run 18, mag. 1000 X.

Figure 23(b). Optical micrograph showing etched surface of the flake obtained from run 20, mag. 1000 X.

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Optical micrograph showing the etched surface of the flake obtained from run 19, mag, 1000 X.

SEM micrograph of etched surface of flake observed in Figure 23(a), mag. 3100 X.

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grain size ingot 970 licrons .

7-cn z o cr o 6 -l-H

LU IVJ I—I CJ1

cr CD

5 30 35, , 40 VELOCITY (m/sl

Figure 25. Variation of grain size with disk velocity.

s

I

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particular alloy composition { Ino.44Gao.56Sb' 1,1111 19 '

Figures 26 and 27 show the bright field TEM image and an

electron diffraction pattern of the same area obtained on a flake

from run 13. The diffraction pattern shows the presence of rings with

spots indicating the flakes to be microcrystalline. The presence of

spots is because of the random orientation of the small crystallites

which intersect the electron beam producing a diffraction pattern with

a series of spots. The bright field image in Fig. 26 shows presence of

several phases. Fig.28 is a diffraction pattern of powder from run 20

(20P). The diffraction pattern shows rings with an amorphous halo,

which indicates the presence of a small fraction of amorphous phase.

However exposure of the sample to the electron beam during microscopy

caused crystallization indicating that the alloy was highly beam

sensitive. Evidence of this is seen in Fig.30 which is the diffraction

pattern of the same area after exposing the sample to the beam for ten

minutes. A distinct sharpness in the rings is seen which indicates

microcrystallinity in the sample. The rings in the diffraction patterns

correspond to a diamond cubic structure. Bright field image of the

same area around the halo in Fig. 28 showed the formation of very small

crystallites of the order of 32 nm. The bright field image is shown in

Fig.29.

Dark field image of the same area obtained by tilting the beam

is shown in Fig. 31. The image shows the tiny crystallites, however

the streaking in the image is because of the instrument drift induced

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t 0-4 Mm |

Figure 26. Bright field image of flake obtained from run 13, mag. 80,000 X.

f

s

I

Figure 27. Diffraction pattern obtained from an area of flake observed in Figure 26.

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8 3

Figure 28. Diffraction pattern obtained for powder obtained from run 20.

Figure 29. Bright field image of the same region from which diffraction pattern in Figure 28 was obtained, mag. 31,200 X.

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8 4

Figure 30. Diffraction pattern obtained for powder obtained from run 20 after exposing the sample to the electron beam for 10 minutes.

, - > ' * ** Y s*. 1

* • / fy S' />;'

/ ;

Figure 31. Dark field image of the region from which diffraction pattern in Figure 28 was obtained, mag. 31,200 X.

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85

in the sample. This has become very prominent due to the very long

exposure of hundred seconds introduced during photography of the image.

EDS spectra of these areas were obtained to confirm that the patterns

were originating from the sample and not from the carbon film on the

grid. The spectra confirmed the presence of In, Ga and Sb.

Auger Electron Spectroscopy performed on the surface of a flake

( obtained from run 14) shows the presence of In, Ga, Sb, 0,C and Cu.

The presence of Cu on the suface is possibly due to quenching on the

copper disk. Fig.32 shows the Auger electron spectra indicating the

presence of the individual elements. Depth profiling was performed to

determine if oxide formation occured in the processing of these alloys.

The profiling was obtained by argon ion beam etching while continuously

monitoring the elemental Auger lines. The ion gun energy used was 2 keV

and the electron beam vised was 3 keV. The depth profiles are shown in

Fig.33 which do not show any substantial amounts of 0, C and Cu in

the sample for a sputtering time of ten minutes.

4.5 Transformations

In order to detect the presence of any metastable phases DSC

analyses were conducted on the ingot and melt spun samples of

composition Ino.44Gao.56s obtained from run 18,19 and 20. The

analyses were conducted on the powders as well as the flakes. These DSC

traces for the ingot and the flakes are shown in Fig.34. The DSC trace

for the ingot shows an exotherm at = 300°C which confirms the non-

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• •

• •

; Gfl C- IN-h

: N 1

SB

cu G fl

>

1111 i

:

: 0 l

:

i

180 200 300 400 S00 600 700 800 900 1000 1100 1200 KINETIC ENERGY, EV

Figure 32. Auger electron spectra showing presence of C, 02, Cu, In, Ga, Sb in flake obtained from run 16.

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J0

a8

6

4

2

0

J0

a8

6

4

2

0

. . :

:: : A :: :

A

UUl A \ I

: n. \ f / \ :

: n. V V J 1 : n.

3 4 5 6 7 8 SPUTTER TIHE (HIH.)

10

T : • . •

; ;

J ;

i

: • : •

1 c 1 i

: •

l i l l l 2 3 4 5 6 7

SPUTTER TIHE (HIH.) 10

i i i i t i i i i i i i i i M i i i i i i i »i i ii ii i

Or 1

2 3 4 5 6 7 8 9 1 0 SPUTTER TIHE (HIN.)

Figure 33. Depth profiles for Cu, C and elemental 0 in flake obtained from run 16.

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Ingot

t -a}

i

t -ioj

i-0 lea 2 ee

f: »• 300 400 S00 800 7£

Tump«raiur<i C#C5

Figure 34. DSC traces of ingot and flakes obtained from run 18,19 and 20.

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89

equilibrium nature of the ingot indicated by the X-ray results. The X-

ray results for the ingot as described earlier indicated presence of

non-equilibrium phases confirmed by the lattice parameter calculations.

Similar exotherms were seen for the flakes melt spun at low and

high rpms corresponding to run 18 and 20 identified as 18A and 20F.

This indicates that the flakes melt spun at relatively low and high

rpm's retain metastable phases identical to the ingot. But the DSC

trace for the flakes obtained from run 19 identified as 19F does not

show the exothermic peak. Hie DSC traces for the corresponding powders

are shown in Fig.35 which show a similar trend to those obtained for

the flakes. The only difference is that the trace identified as 18B and

18C which corresponds to the fines obtained from run 18 in addition to

the powder obtained from run 19 does not exhibit the exothermic peak.

This observation supported by the grain size-velocity plot (see

Fig.25) suggests the possibility of an amorphous phase being formed in

this sample although the X-ray traces were similar to those obtained in

the case of the powders and flakes obtained from runs 18 and 20. (see

Fig. 12). Also, the fraction of the crystalline phase may be

insufficient for the DSC to register the exothemic peak characteristic

to the phase. A summary of these conclusions are shown in Table 11.

Thermal analyses were also conducted on the powders obtained

from run 8,9,10 for samples of composition Ino.51GaO 4gSb. The DSC

traces for these runs are shown in Fig.36. There is no evidence of any

exothermic transition, however the trace for run 10 where the melt was

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0->

20 P

18C

-3

-4-

B 0 U.

18 B 0 1 -3"

- j . .

19 P

-3-

- S -

IBS

Figure 35. DSC traces of corresponding powders obtcdned from runs 18, 19 and 20.

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Table 11. Summary of XRD and DSC Analysis

Sample Composition IriQ. 44Ga0.56s*3

Form Run $

Phase Present (XRD)

Phase present (DSC)

Flake 18A A A

Flake 18B A -

Powder 18C A -

Powder 19P A -

Flake 19F A

Flake 20F A A

Powder 20P A A

Ingot A A

* The letter 'A' represents the metastable phase

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1 0 -

0-

-10..

-10

0

-10

» -20 o u. .*> s -30 X

-40-

-50. .

-ea

run 10

run 9

run 8

100 200 300 400 500 600 700 T«npTatur« <®D

Figure 36. DSC trace of powder obtained from runs 8, 9, and 10.

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held for 18 seconds before ejecting the liquid metal shows a slight but

noticeable change in slope at approximately 400°C which indicates a

probable glass transition. This reasoning can be supported by the

observation of certain unetched regions in the microstructure of this

sample (see Fig. 22(c)). Apart from this observation all the traces

exhibit consistently identical endotherms.

4.6 Electrical Property Measurements

Bulk electrical resistivity measurements were made at room

temperature only on flakes of composition InQ GaQ gSb obtained from

run 7. The flakes were large enough to permit electrical resistivity

measurements using the four point probe. Flakes obtained by melt

spinning other compositions were too small to perform electrical

resistivity measurements. The measurements were made on flakes

annealed at temperatures of 100,200,300,and 400°C for a constant period

of two hours. The resistivity measurements were made by incorporating

corrections for the thickness and the shape of the flake. [97].

A plot of these values for the various temperatures is shown in

Fig. 37. Each resistivity value represents an average of twenty

measurements performed on each flake. The bar around each point is a

measure of standard deviation. The plot is unusual since one would

expect the resistivity to decrease with increasing annealing

temperatures due to stress relaxation. Resistivity values reported at

room temperature by earlier investigators for bulk ingots of similar

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• unquenched ingot 17 .109-

.099-

.079-

.069-

.059-

.049-

.04-

.03-

.02-1

.01-

400 600 Too 300 ANNEALING TEMPERATURE (degree K)

200 800

Figure 37. Variation of bulk electrical resistivity with annealing temperatures for flakes obtained from run

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95

compositions are shown in Table 12. The values obtained by them fall in

the vicinity of the values obtained in the present investigation for

annealing temperatures greater than 200°c.

An explanation for this behaviour is suggested: The rapidly

solidified flakes contain, in part, a metastable phase as inferred from

the X-ray analysis and different metastable phases can have different

resistivity values. Thus the rapidly quenched flake consisting of

metastable phases at room temperature has a resistivity value of 0.02

ohm-cm which is higher than that measured for the ingot. The higher

value of the quenched flake is due to the disorder induced in the

sample by the quenching process[96]. On subsequent annealing there is a

reduction in resistivity due to an ordering induced by the stress

relaxation. Further annealing probably results in the flake

microstructure passing through metastable phases which have

resistivity values in the vicinity of those reported by other

investigators in Table 12. Additional experiments would need to be

performed to confirm this hypothesis.

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Table 12. Resistivity Values cited in Literature

Bulk electrical Composition Temperature Reference Resistivity,

( C) ( ohm-cm )

8.50 x 10~2 In0.5Ga0.5Sb 30 Coderre and Woolley [92].

3.16 x 10 2 la-. CGSQ cSb 30 Omskii and Kblaniets [95].

1.99 x 10~2 InQ>5GaQ 5Sb 30 Omskii and Kblaniets [93].

7.90 x 10~2 Iitq 53Gao 47Sb 30 Woolley and Gillett[94],

1.10 x 10"1 In0.52Ga0.48Sb 30 Joullie et al.[78].

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CHAPTER 5 /

SUMMARY AND CONCLUSIONS

The results obtained during the present investigation allow the

following conclusions:

1. Rapid solidification of several InxGa1_xSb compositions resulted in

the production of flakes or powders depending on the rpm of the copper

disk. Surface examination of the flakes showed presence of second phase

precipitates of the order of 1 micron which were rich in InSb. The

flake thicknesses ranged from 15 microns to 50 microns for varying disk

speeds. The powders obtained were irregular and varied in sizes from

one micron to several hundred microns.

2. Furnace cooled ingots yield a mixture of non equilibrium phases

which are retained in the powders and flakes on subsequent rapid

solidification.

3. X-ray and TEM analysis showed the flakes and powders to be

microcrystalline. The grain sizes varied With the disk velocity

exhibiting a minimum grain size of 2.7 microns in the case of the

flakes produced by rapid solidification in the intermediate velocity

range of 45-60 m/s. The surface of the flakes showed fine cellular

morphology whereas the cross sections of both the powder and the flake

showed presence of fine columnar grains. Powders obtained from rapid

solidification at high rpm showed presence of an amorphous phase which

97

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is highly electron beam sensitive.

4. Thermal analysis indicated suppression of the metastable phase

showed by X-ray analysis in the case of the flakes and powders obtained

by rapid solidification in the intermediate velocity range of 45-60

m/s.

5. The flakes do not show presence of any substantial amounts of oxygen

or copper on the surface.

6. Bulk electrical resistivity measurements on flakes of composition

In0.51Ga0.49sb showed a resistivity minimum at an annealing temperature

of 100°C.

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REFERENCES

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2. Ananthraman, T. R., and C. Suryanarayana. J. Mat. Sc., Vol.6, p. 1111, 1971.

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4. Duwez, P., R. H. Willens, and W. Klement. J. App. Phys. Vol.31, p. 1136, 1960.

5. Ananthraman, T. R. Metallic glasses. Production, properties and application. T. R. Ananthraman, ed., Switzerland, Rockport, Ma: TransTech publications, p. 1, 1984.

6. Guntherodt, H. J. Proceedings of the Fifth International Conference on Rapidly Quenched Metals. S. Steeb and H. Warlimont, eds. Elsevier Science Publishers B. V., p. 591, 1985.

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99

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15. Miroschnichenko, I. S. and V. 0. Zakharov. Industr. Lab., Vol. 35, p. 362, 1969.

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17. Lohberg, K and H. Muller. Fizika 2, Suppl. 2, paper 4, 1970.

18. Brower, W. E. Jr., R. Strachan, and M. C. Flemings. Cast Metals Res. J., Vol. 6, p. 176, 1970.

19. Ruhl, R. C. Mat. Sc. & Eng., Vol. 1, p. 313, 1967.

20. Jones, H. Mat. Sc. & Eng., Vol. 5, p. 297, 1969/1970.

21. Jones, H. Mat. Sc. & Eng., Vol. 5. p. 1, 1969.

22. Toda, T and R. Maddin., Trans. Met. Soc. AIME, Vol. 245, p. 1045, 1969.

23. Chen, H. S. and D. Turnbull. Acta Metall., Vol. 17, p. 1021, 1969.

24. Davies, H. A. and J. B. Hull, Scripta Metall., Vol. 6, p. 241, 1972.

25. Amorphous Glassy Metal and Allays, Tech. Rep. AFML-TR-78-70, 1977.

26. Beghl, G, R. Matera, and G. J. Piatti, Mat. Sc., Vol. 5, p. 820, 1970.

27. Pietrokowsky, P, Rev. Sc. Instrum., Vol. 34, p. 445, 1963.

28. Pond, R and R. Maddin. Trans. Met. Soc. AIME, Vol. 245, p. 2475, 1969.

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