solidification of nickel-base alloys containing … · 2020. 4. 2. · solidification of...

SOLIDIFICATION OF NICKEL-BASE ALLOYSCONTAINING TITANIUM AND ALUMINUM

Item Type text; Dissertation-Reproduction (electronic)

Authors Vaughn, Glen Allen

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 27/08/2021 07:00:42

Link to Item http://hdl.handle.net/10150/298843

http://hdl.handle.net/10150/298843

INFORMATION TO USERS

This material was produced from a microfilm copy of the original document. While the most advanced technological means to photograph and reproduce this document have been used, the quality is heavily dependent upon the quality of the original submitted.

The following explanation of techniques is provided to help you understand markings or patterns which may appear on this reproduction.

1. The sign or "target" for pages apparently lacking from the document photographed is "Missing Page(s)". If it was possible to obtain the missing page(s) or section, they are spliced into the film along with adjacent pages. This may have necessitated cutting thru an image and duplicating adjacent pages to insure you complete continuity.

2. When an image on the film is obliterated with a large round black mark, it is an indication that the photographer suspected that the copy may have moved during exposure and thus cause a blurred image. You will find a

good image of the page in the adjacent frame.

3. When a map, drawing or chart, etc., was part of the material being photographed the photographer followed a definite method in "sectioning" the material. It is customary to begin photoing at the upper left hand corner of a large sheet and to continue photoing from left to right in equal sections with a small overlap. If necessary, sectioning is continued again — beginning below the first row and continuing on until complete.

4. The majority of users indicate that the textual content is of greatest value, however, a somewhat higher quality reproduction could be made from "photographs" if essential to the understanding of the dissertation. Silver prints of "photographs" may be ordered at additional charge by writing the Order Department, giving the catalog number, title, author and specific pages you wish reproduced.

5. PLEASE NOTE: Some pages may have indistinct print. Filmed as received.

University Microfilms International 300 North Zeeb Road Ann Arbor, Michigan 48106 USA

St. John's Road, Tyler's Green High Wycombe, Bucks, England HP10 8HR

7810621

VAUGHN, GLEN ALLEN SOLIDIFICATION OF NJCKEL-BA9E ALLEYS CONTAINING TITANIUM AND ALUMINUM,

THe UNIVERSITY OF ARIZONA# PH.O,# 1978

University Microfilms

International 300 n. zeeb road, ann arbor, mi 4bio6

SOLIDIFICATION OF NICKEL-BASE ALLOYS

CONTAINING TITANIUM AND ALUMINUM

by

Glen Allen Vaughn

A Dissertation Submitted to the Faculty of the

DEPARTMENT OF METALLURGICAL ENGINEERING

In Partial Fulfillment of the Requirements For the Degree of

DOCTOR OF PHILOSOPHY WITH A MAJOR IN METALLURGY

In the Graduate College

THE UNIVERSITY OF ARIZONA

19 7 8

THE UNIVERSITY OF ARIZONA

GRADUATE COLLEGE

I hereby recommend that this dissertation prepared under my

direction by Glen Allen Vaughn

entitled SOLIDIFICATION OF NICKEL-BASE ALLOYS

CONTAINING TITANIUM AND ALUMINUM

be accepted as fulfilling the dissertation requirement for the

degree of Doctor of Philosophy

// „ ( Y l , —

Dissertation Director

7/7/ 7Date

As members of the Final Examination Committee, we certify

that we have read this dissertation and agree that it may be

presented for final defense.

/n rj/

4-/u/7p>

C-/

Final approval and acceptance of this dissertation is contingent on the candidate's adequate performance and defense thereof at the final oral examination.

STATEMENT BY AUTHOR

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

Brief quotations from this dissertation are allowable without special permission, provided that accurate acknowledgment of source is made. Requests for permission for extended quotation from or reproduction of this manuscript in whole or in part may be granted ty the head of the major department or the Dean of the Graduate College when in his judgment 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: V

To the memory of my mother,

Melba LaVerne Browning Vaughn,

whose recent death has left

an irreplaceable void in my life.

iii

ACKNOWLEDGMENTS

The author wishes to express his gratitude to his

advisor, Dr. Gordon H. Geiger, for his assistance and

guidance. The assistance provided by Dr. Kenneth Keating

and Dr. Louis Demer were much appreciated. I wish to

thank Mr. Thomas Teska, for his help when problems

developed with the electron microprobe analyzer. The

financial and experimental assistance provided by Special

Metals Corporation and the advice from its staff members,

especially Dr. Willard Sutton, were greatly appreciated.

Many experiments required an extra hand and my fellow

graduate student, Mr. John Smith, was always there to

help; his assistance and friendship made the work a little

easier. Finally, a special thanks to my wife, Robin, for

typing this manuscript and putting up with me during my

years in graduate school.

iv

TABLE OF CONTENTS

Page

LIST OF ILLUSTRATIONS vii

LIST OF TABLES xi

ABSTRACT xii

1. INTRODUCTION 1

1.1 Purpose of the Investigation 1 1.2 General Use of Nickel-Base

Superalloys 3 1.3 Physical Metallurgy of

Superalloys 3 1.4 Vacuum Induction Melting of

Superalloys 5

2. SOLIDIFICATION THEORY 7

2.1 Homogeneous Nucleation 7 2.2 Heterogeneous Nucleation ....... 10 2.3 Solidification of Pure Metals .... 11 2.4 Solidification of Alloys 13 2.5 Constitutional Supercooling 15 2.6 Dendrite Structure in Alloys 18 2.7 Dendrite Arm Spacing 20 2.8 Microsegregation 22

3. THERMODYNAMICS OF DILUTE LIQUID NICKEL ALLOYS 26

3.1 Deoxidation Equilibria in Liquid Nickel Alloys 27

4. INCLUSIONS FORMED IN CAST STRUCTURES ... 28

4.1 Primary Inclusions 28 4.2 Secondary Inclusions 32 4.3 Significance of Inclusions 33

v

vi

TABLE OF CONTENTS--Continued

Page

5. EXPERIMENTAL PROCEDURE 34

5.1 Preparation of the Alloys 34 5.2 Melting of the Alloys 37 5.3 Molds 40 5.4 Casting of the Alloys 42 5.5 Electron Microprobe Analysis 45 5.6 Metallographic Analysis 49 5.7 Differential Thermal Analysis .... 50

6. RESULTS AND DISCUSSION 52

6.1 Nickel-Aluminum and Nickel-Titanium Phase Diagrams 57

6.2 The Nickel-Titanium-Aluminum System . 64 6.3 Microstructure of the As-Cast

Alloys 69 6.4 Dendritic Microsegregation 89 6.5 Macrosegregation 95 6.6 Variation of Secondary Arm

Spacing with Cooling Rate 97 6.7 Alloy Melt-Crucible Interactions ... 99

7. CONCLUSIONS 139

APPENDIX A: THERMODYNAMIC DATA 142

APPENDIX B: DESCRIPTION OF ZAF PROGRAM . . 164

APPENDIX C: MICROSEGREGATION DATA .... 170

LIST OF REFERENCES 211

LIST OF ILLUSTRATIONS

Figure Page

1 Sequence of events for positive temperature gradient in the liquid and solid 12

2 Sequence of events for negative temperature gradient in the liquid and positive temperature gradient in the solid 14

3 Portion of a phase diagram for an alloy of composition CQ 16

4 Constitutional supercooling ahead of an interface 16

5 Schematic view of dendrite array showing lateral solute transport 23

6 The MgO and AI2O3 crucibles 39

7 Chill and insulating molds 41

8 Melting set-up inside vacuum chamber .... 43

9 Control panel for vacuum melting apparatus 44

10 Scanning Electron Microprobe Quantometer (SEMQ) 46

11 Schematic showing location of SEMQ specimen 47

12 Nickel-aluminum binary system 58

13 Nickel-titanium binary system: Version I . . -59

14 Nickel-titanium binary system: Version II. . 60

15 The y' precipitate in the y matrix 62

16 The n phase precipitate in the y matrix ... 63

17 Isothermal sections for the nickel-titan ium-aluminum phase diagrams at 1150°C and 1000°C 65

vii

viii

LIST OF ILLUSTRATIONS--Continued

Figure Page

18 Isothermal sections for the nickel-titanium-aluminum phase diagram at 850°C and 750°C 66

19 Liquidus surface of the nickel-rich end of the nickel-titanium-aluminum phase diagram 70

20 Vertical portion of the nickel-titanium-aluminum system at 90w/o nickel 71

21 Alloy composition investigated plotted on the liquidus surface of the nickel rich end of the nickel-titanium-aluminum phase diagram 72

22 Microstructure of 90w/o Ni-10w/o A1 . . . . 74

23 Microstructure of 90w/o Ni-10w/o Ti . . . . 77

24 Micros tructure of 90w/o L,:Ii-10w/o Ti showing n-y eutectic between y dendrite arms 78

25 Microstructure of 89.5w/o Ni-8.5w/o Al-2w/o Ti 82

26 Microstructure of 90w/o Ni-5w/o Al-5w/o Ti 83

27 Microstructure of 90w/o Ni-2w/o Al-8w/o Ti 84

28 Microstructure of Y-Y' mixture between dendrite arms 85

29 Microstructure of 93w/o-5w/o Al-2w/o Ti . . 87

30 Macroscopic Scan from the casting edge to casting center for nickel, aluminum and titanium 96

31 Variation of secondary dendrite arm spacing with local cooling rate 98

32 Variation of secondary dendrite arm spacing with local solidification time. . 100

ix


Figure Page

33 Secondary dendrite arm spacing as a function of local solidification time for several alloys 101

34 Equilibrium relationship between dissolved oxygen in the nickel melt and the content of various deoxidizing elements 103

35 Equilibrium partial pressure of Mg in the gas as a function of w/o C in the melt 106

36 Metallostatic head pressure as a function of depth below the top surface of the melt 108

37 Surface of the alloy melt immediately after melting 109

38 Surface of the alloy melt 5 minutes after melting 110

39 Backscattered x-ray image of typical AloOo clusters found in 90w/o Ni-10w/o Al casting 112

40 Distribution image of Ni-Ka radiation of AI2O3 clusters in Figure 39 113

41 Distribution image of Al-Ka radiation of AI2O3 clusters in Figure 39 114

42 Distribution image of 0-Ka radiation of AI2O3 clusters in Figure 39 115

43 Backscattered x-ray image of MgO particle found in an alloy melted in an MgO crucible 116

44 Distribution image of Ni-Ka radiation of particle found in Figure 43 117

45 Distribution image of Al-Ka radiation of particle in Figure 43 118

X


Figure Page

46 Distribution image of 0-Ka radiation of particle in Figure 43 119

.47 Distribution image of Mg-Ka radiation of particle in Figure 43 120

48 Distribution image of Ti-Ka radiation of particle in Figure 43 121

49 Backscattered x-ray image of Al^O^ particle found in an alloy melted in an A^O^ crucible 125

50 Distribution image of Ni-Ka radiation of AI2O3 particle in Figure 49 126

51 Distribution image of Al-Ka radiation of AI2O3 particle in Figure 49 127

52 Distribution image of 0-Ka radiation of AI2O3 particle in Figure 49 128

53 Distribution image of Ti-Ka radiation of AI2O3 particle in Figure 49 129

54 Schematic of a crucible wall cross section showing reduction of wall thickness near the melt-vacuum boundary 131

55 Primary inclusions in dendrites of 93w/o Ni-5w/o Al-2w/o Ti 132

56 The C + 0 = C0/2\ relationship at CO partial pressures of 0.07 and 0.012 atm 136

LIST OF TABLES

Table Page

1 Analysis of Charge Materials 35

2 Chemical Analysis of Crucibles 38

3 Casting Variables 53

4 Spectrographic Analysis of Cast Alloys . . 55

5 Differential Thermal Analysis Data .... 68

6 Percentage of y-y' Eutectic Mixture in the Alloys 80

7 Segregation Ratio Data for Ni-Al-Ti Alloys 90

8 Recommended Minimum Vacuum Pressures for Ni-Base Alloys 138

xi

ABSTRACT

A series of nickel-titanium-aluminum alloys

containing approximately 90w./o Ni and varying amounts of

titanium and aluminum, was induction-melted in a vacuum __ O

of 1 x 10 torr in either MgO or A^O^ crucibles and

vacuum-cast into either a cylindrical chill or insulating

mold. These alloys were investigated to determine their

solidification behavior, to establish the effect of

solidification rate on microstructure, and to determine the

effect of alloy melt-crucible interactions on

inclusion formation.

As a manifestation of nonequilibrium solidification

the structure of these alloys consisted of a y-y' mixture

between the cored y dendritic arms. The segregation of

titanium and aluminum was strongly affected by alloy

composition. For 90w/o Ni-Al-Ti alloys, the segregation

ratios of titanium and aluminum in the dendrite arms were

found to increase from maximum values of 1.68 and 1.20,

respectively to minimum values of 1.09 and 1.08, respectively

as the Al/Ti ratio increased. The nickel-base alloys were

investigated over a local cooling rate range of 0.46°C/sec

to 5.5°C/sec. The cooling rate had little effect on the

microsegregation of aluminum and titanium. The dependence

of secondary arm spacing with local cooling rate was found

xii

to remain fairly constant for the 90w/o Ni-Al-Ti alloys;

-0 ?7 this relationship was established as d = 27.7 (GR)

The majority of the inclusions found in these cast

alloys were primary inclusions of AÔg. These inclusions

developed as a result of the dissociation of the MgO and

crucibles under vacuum melting of these alloys.

The dissociation of these refractories led to oxygen being

pumped into the melt. This oxygen reacted with the

dissolved aluminum in the melt, rather than the dissolved

titanium, since aluminum has a much greater deoxidation

constant than titanium in nickel-base alloy melts. Those

alloys melted in MgO, in addition to having AÔ^ inclusions,

contained particles of MgO which reacted with the dissolved

aluminum and oxygen in the melt to form MgAÔ^. Although

titanium segregated to a greater extent than aluminum in

the interdendritic regions, the aluminum still controlled

the soluble oxygen content in these areas and only AÔg

secondary inclusions formed.

CHAPTER 1

INTRODUCTION

1.1 Purpose of the Investigation

In the past fifteen years, significant advances

have been made in the description and understanding of

solidification in commercially important alloys. The

concepts of solidification theory pertinent to superalloys

have been reviewed, and a need has been shown for the

development of additional solidification data on

superalloys. It is of particular interest to know the

extent of aluminum and titanium segregation in the super-

alloys since these elements form the Ni^(Al,Ti)

precipitate, which provides the superalloy its greatest

amount of strengthening. Microsegregation of aluminum

and titanium can result in a non-uniform precipitation and

a loss in strength. Thus, to gain an understanding of the

microsegregation of aluminum and titanium in the complex

superalloys, solidification studies of nickel-base alloys

containing alloy additions of titanium and aluminum were

initiated. The effect of such parameters as cooling rate

and alloy composition on microsegregation and alloy

microstructure was investigated.

1

2

The prime reason for using superalloys is their

outstanding strength over the temperature range from

1400 to 1900°F. Nonmetallic inclusions can lower the

mechanical properties of these alloys. Therefore, it is

of extreme importance to eliminate inclusions in

superalloys. Some superalloys are vacuum-induction

melted and vacuum cast to achieve a low dissolved oxygen

content and a subsequent low volume fraction of

inclusions. However, vacuum-induction melted superalloys

have not achieved the low oxygen levels as predicted by

thermodynamic calculations. One possible reason the

volume fraction of inclusions exceeds the predicted value

is a consequence of nonequilibrium cooling conditions. As

a result of nonequilibrium solidification, solutes are

rejected to the interdendritic regions; if the oxygen

and metal solutes (such as titaniTjm and aluminum) reach a

sufficient concentration, an interaction will occur and

inclusions are formed. It is also possible that inclusions

result from melt-crucible interactions. Since inclusions

may form from oxygen dissolved in the melt upon solidifica

tion, or result from particles of insoluble refractory

removed from the crucible, the reactions between nickel-

base alloy melts and crucibles of MgO and were

studied in order to determine how best to minimize oxide

inclusion. These results were verified using thermodynamic

calculations.

3

The knowledge gained from this investigation of

nickel-base alloys containing titanium and aluminum

additions will be used to provide a greater understanding

of the solidification and melt-crucible interaction

behavior of the more complex superalloys,

1.2 General Use of Nickel-Base Superalloys

The need for more heat resistant materials in

aircraft engines led to the development of nickel-base

superalloys. The largest use for nickel-base superalloys

is in the gas turbine industry. Contemporary engines use

these alloys for turbine blades, wheels, shafts, and vanes.

These parts are subjected to high temperatures and cyclic

operations; thermal gradients from heating and cooling

turbine sections of varying size induce thermal stresses,

which subject the blade airfoils to thermal fatigue.

Therefore, superalloys used for these applications must

have good fatigue and creep resistance.

In addition to the gas turbine industry, these

alloys are used in nuclear reactors, furnaces, and a

number of highly specialized products.

1.3 Physical Metallurgy of Superalloys

Nickel-base superalloys have austenitic structures.

Strengthening of the face centered cubic (fee) austenitic

matrix generally falls into the category of solid solution

4

strengthening. To some extent, every addition to the

nickel-base serves as a solid solution strengthening

agent. The solid solution elements typically found in

the austenitic, y, matrix are likely to include iron,

chromium, tungsten, cobalt and molybdenum. The difference

in atomic diameter from that of nickel varies from 41?0 for

cobalt to +137o for tungsten. These elements produce

localized elastic strain fields in the y matrix and these

fields interact with those of dislocations, thereby

increasing the strength of the matrix.

The greatest amount of strengthening in superalloys,

however, is developed by precipitation hardening.

Titanium and aluminum are the most important elements

added to these alloys to make them precipitation

hardenable. The heat treatment consists of solution

treatment followed by aging. Aging produces a dispersion

of a coherent, stable, intermetallic compound named gamma

prime, y'. The basis of gamma prime is the intermetallic

compound Ni^Al. Aluminum may be considered to be the

primary gamma prime forming constituent in the nickel-

base superalloys. Substituting titanium for aluminum

changes the gamma prime morphology from cubic to

spheroidal. The strength of a given alloy is dependent

upon such factors as volume fraction, particle size,

coarsening rate, and composition of the y' precipitate'.

5

All of these factors can be controlled to varying degrees

by heat treatment. Coherency strains and disregistry

between the crystal lattices of the austenite matrix (y)

and the gamma prime, y1, precipitates have been used to

explain the hardening of y'-strengthened superalloys by

Bigelow (1) and by Mihalisin and Decker (2).

1.4 Vacuum Induction Melting of Superalloys

Vacuum induction melting of superalloys was

introduced in the 1950s by the Kelsey-Hayes Company (later

Special Metals Company). Vacuum induction melting (VIM)

will prevent the solution of gases in the melt and remove

dissolved gases from the melt. However, VIM itself will

not completely eliminate dissolved oxygen in the superalloy.

The aluminum and titanium in the superalloy have a great

affinity for oxygen and nitrogen; therefore, by decreasing

these gas contents in the metal, the number of inclusions

in the alloy is decreased. It is generally agreed that

fatigue resistance can be increased by decreasing the

number of inclusions.

Vacuum induction melting of superalloys leads to

improvement in the mechanical properties of the alloys when

tested at high temperatures. The rupture and creep strength

can be significantly increased as full advantage is taken of

the strengthening effects of aluminum and titanium additions

when they are not tied up as oxides and nitrides.

6

Stegman, Shahiman and Achter (3) have shown that creep

resistance of nickel increases when the nickel is

melted and cast in a vacuum, as compared to nickel

melted and cast in air.

Another advantage of VIM is the high degree of

uniformity in properties of the product from ingot to

ingot and heat to heat. Some of this is due to the high

quality of raw materials and the extreme care used in

production, but most of it is due to closer control of

aluminum and titanium, as compared to the classical air

melting procedures.

CHAPTER 2

SOLIDIFICATION THEORY

Solidification is describable by two rate

parameters: one for the nucleation of the solid phase from

a supersaturated liquid and the other for the growth

process. In this chapter, solidification is described

in terms of both nucleation and growth of the solid. In

addition to these parameters, the structural implications

of the theory of solidification are discussed. The

structural features of interest are the inhomogeneities

produced by chemical segregation, and the size and shapes

of the grains produced in a casting.

2.1 Homogeneous Nucleation

Nucleation may be defined as the formation of a

new phase in a distinct region, separated from the

surroundings by a definite boundary. Homogeneous nucleation

is the formation of a new phase without the help of

impurities or external surfaces. Impurity particles and

external surfaces are taken into account in section 2.2.

The classical theory of nucleation was developed

by Volmer and Weber (4) and Becker and Doring (5) for the

condensation of a pure vapor to form a liquid. According

to the classical nucleation model, embryos of solid-like

7

8

molecules continually form in the liquid by statistical

fluctuation. Each molecule in an embryo is assumed to have

the same free energy as if it were part of a bulk solid at

that temperature. This assumes that the free energy is

independent of size and morphology. Secondly, according

to the classical model, the excess surface energy per unit

area of embryo is the same as that of a macroscopic solid-

liquid interface. Embodied in both of these assumptions is

the basic assumption that there is a discontinuous inter

face between the embryo and the liquid.

When a spherical embryo of solid is formed within

a uniform liquid there will be a change in free energy

associated with the difference in volume free energy of

the atoms in the solid and the liquid. In addition, there

will be a term introduced because a number of the atoms

occur in the transition region between liquid and solid.

These atoms will be in a high energy state and are the

origin of the surface free energy of the embryo.

For a spherical embryo of radius, r, the overall

change in free energy, AG, is given by

AG = ^•1Tr2YLC + (j)Tr3AGv (1)

where is the surface free energy and AG^ is the volume

free energy change. Above the melting point, is

positive and below it AG^ is negative. At some critical

radius, r' , AG is a maximum. Differentiating equation

1 and allowing for negative AGV gives the critical radius:

r* - (2)

If the embryo forms in such a way that its radius,

r, is greater than r*, then this leads to a decrease in

AG with further increase in r. Thus, any particle larger

than r will be a nucleus for growth and any particle •X.

smaller than r" will tend to disappear, since the tendency

must always be to decrease AG.

The surface energy term does not change signifi

cantly with temperature. However, the volume free energy

varies with temperature, becoming larger at low tempera

tures. Thus the critical radius decreases with falling

temperature. At temperatures just below the melting •J..

point, r" is large. The rate of homogeneous nucleation

is, therefore, small at the melting point. As the

temperature is lowered, the critical radius rapidly

decreases in size. Homogeneous nucleation is thus made

easier as the amount of supercooling is increased.

Homogejieous nucleation is difficult to study

because it is not easy to prepare a metal in such a way

that foreign particles have been removed. Turnbull (6)

has accomplished this by subdividing the metal into small

shapes that are isolated from each other. Since there is

a limited number of impurity particles present in the bulk

10

liquid, some of the small drops will not contain foreign

matter and nucleation must be homogeneous. It was found

that the amount of supercooling required for homogeneous

nucleation is very large, approximately 0.2Tm, where Tm

is the absolute melting temperature. Walker (7) has

shown that melts of small quantities of nickel can be

supercooled to 296 below the freezing point. Such super

cooling is never observed in commercial practice;

supercooling usually varies between one and ten degrees

centigrade. Nucleation, then, cannot be homogeneous but

must be heterogeneous.

2.2 Heterogeneous Nucleation

Most actual castings nucleate at much less

supercooling than the maximum observed in the small drop

experiments. This discrepancy is attributed to the

presence of a suitable surface in contact with the liquid.

The nucleation is considered to be heterogeneous and to

take place on the surface of the container or on particles

present in the system. Heterogeneous nucleation can occur

provided some preferential sites exist.

Heterogeneous nucleation theory has been developed

by Turnbull (8), Volmer (9) and more recently, Sundquist

and Mondolfo (10) . They established that the energy

fluctuation required for heterogeneous nucleation, AG, is

11

much less than that for homogeneous nucleation. The

basic reason is that if the new phase finds an impurity

particle to grow upon, it can in effect adopt the relatively

large radius of the particle as its own. This means that

only a slight degree of supercooling is needed in compari

son with that needed for homogeneous nucleation.

2.3 Solidification of Pure Metals

Once nucleation has occurred, crystal growth of

the pure metal begins and the structure that develops can

be related to the growth conditions. For growth of the

interface to occur, the temperature of the interface must

be slightly below the equilibrium freezing temperature.

This means that some supercooling must exist if the

interface is to advance.

Consider a case where the area of solid and liquid

adjacent to the interface shows a positive temperature

gradient in the liquid and solid and the temperature

gradient in the solid is steeper than the gradient in the

liquid because of the higher thermal conductivity of the

solid. In this case, the formation of an unstable

protuberance will melt because the local tip temperature

exceeds the melting temperature. This sequence of events

is illustrated in Figure 1.

If the area of the solid and liquid adjacent to

the interface shows a negative temperature gradient in the

12

SOLID LIQUID

G, positive

Temperature

Gs positive

Distance

Interface — T~

SOLID LIQUIO

Tlocal > 7"rr

SOLID

SOLID / LIQUID

LIQUID

Initial Interface form of interface with shape instability

Final form of interface

Figure 1: Sequence of events for positive temperature gradient in the liquid and solid.

13

liquid and a positive temperature gradient in the solid,

a protuberance on the interface will project into a

region where the local tip temperature is below the

melting temperature. In this case, the protuberance will

grow and the interface thus degenerates and grows

dendritically as shown in Figure 2. Dendritic structures

will be discussed in greater detail in section 2.6.

2.4 Solidification of Alloys

When an alloy solidifies, the solid that forms

generally has a different composition than the liquid

from which it is freezing. Therefore, the distribution of

a solute in the solid will generally be different than it

was in the liquid prior to freezing. This redistribution

of solute produced by solidification is termed segregation

or coring.

Both in pure metals and in alloys the structure

can be directly related to supercooling. In pure metals,

supercooling may be produced only by thermal means. In

alloys, supercooling may be produced indirectly by changes

in temperature and composition. If it is produced by

changes in composition combined with temperature changes

it is referred to as constitutional supercooling, and it

is this type of supercooling that determines the growth

structures found in alloys.

14

SOLID LIQUID

Gs positive

ffL negative

Temperature

Interface

~ Tm

Distance

"loc al < 7"m

SOLID LIQUID

Initial form of interface

SOLID! LIQUID

Interface with shape instability

[UQUID LN SOLID~=> r—nr*

I SOLIDI LIQUID

Subsequent forms after shape instability has grown

Figure 2: Sequence of events for negative temperature gradient in the liquid and positive temperature gradient in the solid.

15

2.5 Constitutional Supercooling

Under non-equilibrium conditions, concentration

gradients develop in the liquid ahead of the solid-liquid

interface because the composition of the forming solid is

different from the composition of the liquid from which it

is freezing. If the concentration of solute in the solid

is less than that of the liquid from which it is forming

there must be a rejection of solute into the liquid at

the solid-liquid interface. If sufficient time is not

allowed for this solute to distribute itself throughout

the remainder of the liquid, a concentration gradient

will develop in the liquid ahead of the interface. This

concentration gradient promotes the constitutional super

cooling that is responsible for the'structure found in

alloy castings.

With the aid of a portion of a phase diagram

shown in Figure 3, constitutional supercooling is

illustrated in Figure 4 for an alloy of concentration CQ.

Figure 4 shows two curves; one curve is a plot of the

actual temperature of the liquid as a function of

distance from the solid interface and the second curve

shows the equilibrium liquidus temperature of the alloy

as a function of distance from the interface. The actual

temperature of the liquid is assumed to rise linearly from

the interface. The equilibrium liquidus temperature varies

16

LIQUID

D O w. a> CL E a> H-

SOLID

Composit ion

Figure 3: Portion of a phase diagram for an alloy of composition C0.

Xm posed temperature Equilibrium graaiems hquiaus

temperature

Constitutionally supercooled zone

Distance ahead of the interface

Figure 4: Constitutional supercooling ahead of an interface.

17

with distance because the lower solute content, the higher

the liquidus temperature; this is seen from the phase

diagram in Figure 3. At the interface, the freezing temper

ature is T-^, but on moving away from the interface, it at

first rises rapidly and then levels off to the temperature,

T2, the temperature at which the bulk composition of the

liquid, CQ, will begin to freeze. Local equilibrium is

assumed to exist at the solid-liquid interface; therefore,

both curves must pass through T^. The two curves also

intersect at a distance x from the interface. Within the

distance x, the liquid lies at a temperature which is below

its freezing point; this region is supercooled.

In describing the manner in which the solute

partitions itself between the liquid and solid phases at

equilibrium conditions, it is convenient to define an

equilibrium ratio k. The partition ratio, k, is defined

as the ratio of the solute concentration in the freezing

solid, Cs, to the solute concentration in the liquid at

the same temperature. If the effect of the solute is to

lower the liquidus temperature, then k < 1. If the effect

of the solute is to raise the liquidus temperature, then

k > 1.

Since solidification of alloys is usually a non-

equilibrium solidification the equilibrium lever rule can not

be used to solve for the fraction solidified at a given

temperature. Gulliver (11), Scheil (12), and Pfann (13)

18

have derived an equation, known as the Scheil equation,

which can be used to determine fraction solidified at a

given temperature under non-equilibrium conditions. The

equations are given as

Cs = k Co(l-£s)(k"1:) (3)

CL * Co

where Cg and is the solute composition of the solid and

liquid respectively, k is the equilibrium partition ratio,

Cq is the initial composition of alloy, and f and f^ is the

fraction solid and fraction liquid respectively.

In the derivation of the Scheil equations, the

following assumptions were made:

(1) Local equilibrium exists at the solid-liquid

interface.

(2) No solid state diffusion.

(3) Complete mixing in the liquid.

In a binary eutectic system with terminal solid

solubility, the Scheil equation predicts for alloys of

constant k that some eutectic will form even if the initial

composition is the terminal solid solution region.

2.6 Dendrite Structure in Alloys

If pure metals freeze under a negative temperature

gradient the flat interface becomes unstable and forms

dendrites as shown in Figure 2. In pure metals supercooling

19

ahead of the interface is obtained only if the real

temperature has a negative gradient so that it falls below

the constant freezing temperature. However, in alloys the

freezing temperature is not constant, but rather it is a

function of composition as given by the liquidus line on

the phase diagram. Hence, in alloys we may obtain super

cooling with a positive temperature gradient. If there is

only minor supercooling, certain preferred regions of the

interface will protrude as spikes into the supercooled

region and once started, will grow more rapidly than

neighboring regions. This develops a structure which is

composed of parallel elements of rod-like form which

normally run in the direction of freezing. These rods are

hexagonal in cross section and appear as an array of

hexagonal cells.

As the conditions promoting constitutional super

cooling become severe, i.e., shallower temperature gradients,

faster growth rates, or higher solute concentrations, the

length of constitutionally supercooled liquid ahead of the

original planar interfaces increases, and the extension of

the cell boundaries increases. Eventually, the extended

cells break down and side branch. The resultant tree-like

structure, called dendritic, is the primary mode of alloy

solidification.

20

Each dendrite comprises a single grain. Secondary

dendrite arms grow from the more prominent primary arms.

Tertiary arms grow from secondary arms. Dendritic growth

is strongly crystallographic. The primary arms and side

branches (secondary and tertiary arms) have their arms

parallel to specific crystallographic directions. In

nickel-base superalloys, which are f.c.c., this is the

<001> direction.

2.7 Dendrite Arm Spacing

Dendrite arm spacing in a given alloy is found to

depend strongly and solely on cooling rate. Flemings (14)

found that the relationship between dendrite arm spacing

and thermal variable has the form:

d = a t| = b(GR)"n (5)

where the exponent n is in the range of 1/3 to 1/2 for

secondary spacing and generally very close to 1/2 for

primary spacing. GR is the cooling rate with the units of or /sec, d is the dendrite arm spacing, and t^ is the local

solidification time. Local solidification time is the time

required for a given fixed location to go from the liquidus

temperature to the solidus temperature for that local

composition.

It is generally found that as one decreases the

grain size the strength of a metal increases. There is a

21

well-known relation called the Petch (15) equation that

shows that strength is proportional to the reciprocal of

the square root of the grain diameter. For cast metals,

however, it is always true that strength improves with

decreasing grain size as demonstrated by Wallace (16).

Many examples are found in the literature, such as the

work done by Passmore, Flemings, and Taylor (17), which

show that dendrite arm spacing of cast structures usually

correlates better with mechanical properties than does

grain size. The work of Frederick and Baily (18) shows

that as the dentrite arm spacing is reduced by increasing

solidification rate, the tensile strength and ductility

of aluminum alloys increase. However, it was found that

the yield strength was not significantly altered by

decreasing the dendrite arm spacing.

In addition, fine dendrite arm spacings in cast

alloys are desirable, since it has been shown by Singh,

Bardes, and Flemings (19) that the homogenization time for

an alloy with nonequilibrium solute segregation is

proportional to the square of the dendrite arm spacing

divided by the diffusion coefficient for diffusion in the

solid. Thus, for a solid-state diffusion coefficient of

-10 2 10 cm /sec, a structure segregated on a 1 micron scale

can be homogenized in a time the order of 100 sec, whereas

a similar structure segregated on a 0.1 millimeter scale

would require 10 days.

22

2.8 Microsegregation

Two kinds of chemical segregation are usually

distinguished. The chemical inhomogeneity occurring over

the distance of dendrite spacings is termed microsegrega-

tion. Chemical inhomogeneity occurring over the distance

of the mold wall to casting center is termed macrosegre-

gation. Only microsegregation will be discussed here.

In dendritic solidification, the solute is virtually

all rejected in the lateral direction into the interden-

dritic liquid as shown in Figure 5. This results in a

variation in the solute concentration between the center

and the outside of a dendrite arm. In extreme cases, the

accumulation of solute between the growing dendrite arms

can lead to the formation of second phases in the inter-

dendritic region in amounts significantly greater than

those predicted from the equilibrium diagram. This type of

segregation is termed microsegregation because it extends

over a length on the order of one-half the dendrite

spacing. One way to characterize the amount of micro-

segregation is to determine the volume fraction of such a

nonequilibrium second phase. However, in many alloys a

second phase does not form, even though it is predicted

by the Scheil equation. The second phase may not form if

the equilibrium partition ratio is not constant or if

sufficient solid state diffusion occurs. Consequently, a

common method of characterizing the amount of

23

Liquid Solid

Lateral Z1

solute transport

Figure 5 : Schematic view of dendrite array showing lateral solute transport.

24

microsegregation is to measure the segregation ratio, SR,

defined as

SR = max. concentration (interdendritic region) min. concentration (dendrite stalks)

There is little information concerning the

segregation in superalloys. However, many observations

have been recorded in steel castings. Weinberg and Buhr

(20) showed that the SR of phosphorous in the dendrite

primaries of a 4340 steel increased from 1.1 to 1.8 as the

primary spacing increased from the chill wall to the center

of the casting. These workers measured no change in SR of

nickel, chromium or manganese with arm spacing. Doherty

and Melford (21) found the SR of chromium in a 0.57oC steel

increased from 3 to 29 as the cooling rate decreased from

2000°C/min at the chill mold wall to 6°C/min at the ingot

center. Flemings (22) found the SR of nickel in a Fe-10%

Ni alloy to increase from 1.32 to 1.38 from the chill wall

to casting center. However, Ahearne and Quigley (23) found

no change in SR of several solutes in a high strength steel

with distance from the chill mold wall.

Addition of a third element to a binary alloy often

significantly affects segregation of the original solute.

An interesting example is the effect of carbon on

segregation of chromium in a Fe-1.5% Cr steel as observed

by Flemings (22). The binary alloy showed no segregation

of chromium but additions of carbon increases the SR of

25

chromium up to as much as 5% at about 1.5% C. In this

same investigation, it was shown that carbon did not have

such a dramatic effect on microsegregation in all iron base

alloys. For example, carbon did not affect microsegrega

tion in a Fe-257o Ni alloy. In a different study,

Kohn (24) found that the SR of nickel increased from 1 to

1.8 as arsenic increased from 0 to 0.127o in an alloy steel.

The SR may be determined with the use of an electron

microprobe. To observe segregation between primary dendrite

arms, sectioning must take place normal to the growth

direction; to observe segregation between secondary dendrite

arms, a plane parallel to the growth direction must be

examined.

CHAPTER 3

THERMODYNAMICS OF DILUTE LIQUID NICKEL ALLOYS

Thermodynamics makes it possible to determine with

certainty what reactions can or cannot happen. The

thermodynamics of dilute nickel alloys has been reviewed

and will be presented below. Knowing the thermodynamics

of dilute nickel alloys makes it possible to determine what

reactions can occur in these alloys. Thermodynamics will

be used in later chapters to predict what type of inclusions

may form before or during solidification.

Thermodynamics may be used to calculate the end

product that a system will reach if it is allowed- to go to

equilibrium. In a complicated nonequilibrium system such

as is often encountered in solidification, thermodynamics

would be almost totally useless if applied to the entire

casting. Therefore, it is a common assumption to assume

that thermodynamics can be applied locally, as discussed by

Darken and Gurry (25). For example in section 2.5, when

discussing the Scheil equation, the assumption of local

equilibrium existing at the dendrite-liquid interface was

made.

26

27

3.1 Deoxidation Equilibria in Liquid Nickel Alloys

A good deal of information on the thermodynamic

behavior of elements in liquid nickel has been reported

in the literature. Unfortunately, these data are widely

scattered and presented in a variety of ways. Some systems

have been reviewed by Hultgren and co-workers (26) in their

survey of the thermodynamic properties of binary metallic

alloys. There is also thermodynamic information in the

surveys on binary phase diagrams by Hansen (27), Elliott

(28), and Shunk (29), but no single compilation has been

made. Consequently, as part of the solidification study

of nickel-base alloys, the available thermodynamic data

was reviewed for binary and ternary alloys and summarized

and presented in a paper by Sigworth, Elliott, Vaughn and

Geiger (30). This paper is presented in Appendix A.

CHAPTER 4

INCLUSIONS FORMED IN CAST STRUCTURES

All cast metals have some inclusions. Inclusions

can be classified in various ways but for the purpose of

this discussion the terms primary and secondary inclusions

shall be used. Primary inclusions are defined as those

which form prior to the solidification of the major

metallic phase; whereas, secondary inclusions form during

or after formation of the major phase.

4.1 Primary Inclusions

Primary inclusions can result when a grain refiner

or a deoxidizer is added to the molten metal. For example

in steel, the deoxidation product Al^O^ forms when

aluminum is added as the deoxidizing agent. This leads to

the formation of AI2O3 inclusions. In vacuum induction

melting and vacuum casting of nickel-base superalloys,

deoxidizers or grain refining agents are not added.

However, it is still possible that primary inclusions can

form in these vacuum melted and cast nickel alloys. It

can be shown by using the thermodynamic relationships

presented in Appendix A, that when a molten Ni-lw/o Al

alloy at 1500°C contains as little as 0.0000065 w/o 0,

AI2O3 should form. Thus, there is a possibility if a. source

28

29

of oxygen exists, that primary inclusions will form

in nickel alloys containing aluminum as an alloying

constituent. It is known that the kinetics of the reaction

3 A1 + 20 = AÔg, is very fast in liquid iron alloys (31);

therefore, it can be assumed that the kinetics of this

reaction is similar in liquid nickel alloys. In this case,

AÔg will form as primary inclusions during vacuum melting

of nickel alloys containing aluminum.

In liquid steels deoxidized with aluminum, it was

noted by To.rsell ana Olette (32) that primary AÔg inclu

sions are initially only a few microns in size but increase

in size with time. Elliott, Iguchi, and Chiang (33) have

observed that AI2O3 inclusions collide in the liquid steel

but the particles do not coalesce. The result is that

large interconnected clusters form, which contained a

hundred or more individual inclusions.

Non-metallic particles may enter the cast structure

from outside sources such as the refractory crucible, runner,

pouring spout or slag. Such inclusions are usually referred

to as exogenous. In addition, inclusions may result from

crucible-melt reactions. The chemical reactions that occur

between vacuum melted alloys and their crucibles have been

studied by a number of investigators.

30

Olen, Gonano, and Heck (34) studied a series of

17o C-Fe melts under various partial pressures of carbon

monoxide in a 9970 A^O^ crucible at 1580°c. Vacuum fusion

analysis of the samples indicated that the oxygen contents

were not affected by processing the melts at pressures less

than 100 torr. It was hypothesized that melt-refractory

interactions limited the minimum oxygen content attainable

during the vacuum-carbon deoxidation. Oberg et al.(35)

studied the variations in oxygen and carbon content during

vacuum deoxidation of steel. In this study, a steel with

an initial carbon content of 1% was melted in a MgO crucible

under a vacuum of 10"^ torr. All the carbon was consumed

during the carbon boil after three hours. The oxygen con

centration decreased quickly during the boil from 40 to 4

ppm. At the end of the boil, the oxygen content of the melt

increased with time. One hour after the carbon boil ended,

the oxygen content increased from 4 to 148 ppm and after

another hour increased to 483 ppm. It was assumed that

oxygen was continuously brought into the liquid steel as a

result of the dissolution of the MgO crucible.

Snape and Beeley (36) investigated the refractory-

melt reactions in vacuum induction melted nickel-base alloys.

The reactivity in vacuum of three nickel-base alloys with

alumina (A^O^) , magnesia (MgO) , zirconia (ZrO£) , and

thoria (ThO£) refractories was investigated. The nickel

alloys studied were G-64 (0.12 w/o C, 0.2 w/o Ti, 0.07 w/o

Zr, 5.8W/o Al, 10.5W/o Cr, 0.2W/o Co, 0.4W/o Si, 2.9W/o Mo,

3.9w/o W, bal. Ni), In-100 (0.2w/o C, 4.5w/o Ti, 0.09w/o Zr,

5.1W/o Al, 10.1W/o Cr, 15.0w/o Ci, .05w/o Si, 3.1w/o Mo,

0.65W/o V, bal. Ni), and NiC. The G-64 and In-100 alloys

initially contained 0.0015 w/o 0. When alloy G-14 was

melted in A^O^ and ThC^ crucibles, the oxygen content of

the alloy decreased to 0.0011 w/o and remained constant

after holding the melt in the crucible for 40 minutes.

During this time the carbon content decreased from 0.120 w/o

to 0.115 w/o C. Melting this alloy in MgO and ZrC>2

crucibles caused the oxygen content of the alloy to decrease

to 0.0022 w/o during the first 30 minutes of holding but

the oxygen content then increased to 0.0060 w/o after 60

minutes. The carbon content decreased from 0.12 w/o to

0.105 w/o. Similar results were obtained for the Ni-C and

In-100 alloys. It was assumed that reactions between the

refractories and nickel alloys containing carbon occured by

penetration and solution of the refractory by the metal.

Although the authors discuss the use of dense and porous

crucibles in regards to liquid metal penetration, no density

data for the crucibles used in this investigation was

presented. The authors also assumed that in the case of

dense MgO and ZrC>2 crucibles the reaction represented by the

equation, MemOn + nC = mMe + nCO, occurred. However, no

thermodynamic calculations were made by the investigators to

32

determine if such reactions are possible. Snape and Beeley

did not consider that a crucible such as MgO, containing a

nickel alloy in a vacuum, may dissociate according to MgO =

Mg(g) + 0; therefore, transferring oxygen to the melt. The

work of these investigators, therefore, seems highly

speculative and it appears a better analysis is needed.

Since primary inclusions form before the major

metallic phase, it was once thought that these inclusions

formed within dendrites. However, it is now recognized by

Flemings (14) that primary inclusions can be

pushed by thickening dendrites and thus these inclusions

may appear predominantly in interdendritic spaces.

4.2 Secondary Inclusions

Secondary inclusions result because alloy or

impurity elements are usually rejected to the interdendritic

spaces during solidification. Solutes that lower the

melting point of the alloy, in this case a nickel-base

alloy, are said to have equilibrium partition ratios less

than one. Aluminum and titanium both lower the melting

point of nickel. During the solidification of a nickel

alloy containing aluminum and titanium, these solutes

segregate to regions between the primary and secondary arms.

In addition to these major elements, segregation of trace

impurities such as oxygen also occurs between arms. If the

33

oxygen and metal solutes being rejected to the interden-

dritic region reach sufficient concentration in that region

to cause supersaturation with respect to a thermodynamically

stable oxide, an interaction will occur and a second phase

or secondary inclusion is formed.

Secondary inclusions are usually small compared with

dendrite arm spacing. Flemings (14) has observed that

secondary inclusions are usually in the range of 0.1 to 5

microns for typical ferrous castings. The size and

morphology of the inclusions are dependent upon the

composition and solidification rate of the alloy.

4.3 Significance of Inclusions

The importance of nonmetallic inclusions is in their

ability to affect the mechanical properties of alloys.

Inclusions tend to lower the mechanical properties of alloys.

Wallace (37) points out that the fatigue properties, tensile

ductility, tensile strength and impact ductility decrease as

the number of inclusions increase. The size of inclusions,

as well as the total quantity of inclusions present is also

important. For example, Cummings, Stulen, and Shulte (38)

in an extensive study of ASE 4340 steels found that the

larger an inclusion the more potent it is in starting a

crack. Primary inclusions are usually much larger than those

of secondary inclusions; therefore, primary inclusions are

far more serious, per inclusion, than the latter.

CHAPTER 5

EXPERIMENTAL PROCEDURE

A series of nickel solid solution alloys containing

titanium and aluminum were vacuum induction melted and

vacuum cast into ingots which were 1% inches in diameter.

The nickel content of each alloy was maintained at

approximately 90 w/o while varying the titanium and aluminum.

Initially all the alloys were melted in an MgO crucible.

The MgO crucible was then removed and replaced with an

A^O^ crucible. Alloys identical in composition to those

melted in MgO were melted in A^O^. Both crucibles were

later examined for metal penetration and erosion. After the

alloys were cast, they were sectioned and examined by stan

dard metallographic techniques and with a scanning electron

microprobe quantometer (SEMQ) to determine the type of

structure produced and to evaluate inclusion formation.

5,1 Preparation of the Alloys

The alloys were prepared from electrolytic nickel

and commercially pure aluminum and titanium. The chemical

analysis for each of these constituents is given in Table 1.

For each melt, a 1400 gram sample was prepared.

A series of nickel-titanium-aluminum alloys was

prepared where the composition of nickel was maintained at

34

TABLE 1

ANALYSIS OF CHARGE MATERIALS

Titanium

Ti 99.88W/o

C ,03W/o

Fe .07w/o

Si .02w/o

Nickel

Ni : 99.988w/o

C : .002w/o

Co : .010w/o

Aluminum

A1 : 99,79w/o

C : .01w/o

Fe : .15w/o

Si : .05w/o

36

90w/o while varying the composition of the aluminum and

titanium. To obtain the desired alloy composition, each

component was weighed to the nearest one milligram. After

each alloy was melted and cast, the composition was verified

by spectrographic analysis. In each ingot, a cylindrical

section was removed such that the sample surface chosen for

spectrographic analysis was immediately adjacent to that

part of the casting which was selected for dendrite arm and

microsegregation evaluation. (See section 5.5 for further

discussion.) To obtain a spectrum, an arc was struck on

the specimen to be analyzed producing a burn area approxi

mately %-inch in diameter. Two such burns were made on each

sample analyzed; one was near the outer edge (near chill

wall for mild steel mold) and the other was at the center.

The analyses were performed by Special Metals Corporation

on a three-meter Jerrell-Ash spectrograph. The spectrograms

were recorded on glass photographic plates and read through

a densitometer against standards prepared by Special Metals.

In addition to the spectrographic analysis, an

oxygen and nitrogen determination was made on each sample.

This analysis was performed in a Leco T-30 apparatus by

Special Metals Corporation,

37

5.2 Melting of the Alloys

Initially all the alloys were melted in an MgO

crucible. The MgO crucible was then removed and replaced

with an AÔ^ crucible. Alloys identical in composition

to those melted in MgO were then melted in the AÔ^

crucible. Both the MgO and AÔ^ crucibles were obtained

from the Norton Company. These crucibles had an approximate

nickel alloy melt capacity of five pounds. Dimensions of

both crucibles were as follows: a wall thickness of 5/16

inches, a height of 6 inches, and an inside diameter of

2 3/4 inches. The MgO and AÔ^ crucibles had a nominal

purity of 99w/o. The chemical analyses which were furnished

by Norton for these crucibles is given in Table 2. The

apparent porosity was 22% for the MgO crucible and 19% for

the AÔ^ crucible. These crucibles are shown in Figure 6.

The crucible was centered within the induction coils

of a Stokes Vacuum Casting Unit. The components for a given

alloy were placed in the crucible and the system was

-3 evacuated to 1 x 10 torr. An Inductotherm 15 kilowatt

generator was used to supply power to the coils,

Initially the samples were heated slowly with 5

kilowatts being applied; after 5 minutes the power was

gradually increased to approximately 10 kilowatts. After

all the components had melted, the liquid metal was held

in the refractory crucible for five minutes at a temperature

TABLE 2

CHEMICAL ANALYSIS OF CRUCIBLES

MgO Crucible

MgO - 99.00w/o

CaO - 0.10w/o

Si02 - 0 .50w/o

A1203- 0.35w/O

Fe203- 0.10w/o

A1203 Crucible

A12°3" 99 .01W/o

sio2 - 0 .58w/o

Fe2°3" 0 .llw/o

Na20 - 0 .17w/o

Figure 6: The MgO and Al^Oo crucibles crucible is on the right.

40

near 1500°C prior to pouring into the mold. The temperature

of each melt was monitored by using a Leeds and Northrup

optical pyrometer. The pyrometer used in this casting

unit was calibrated against the melting points of nickel

and iron and was found to read within + 15°C.

5.3 Molds

To establish the relation between dendrite arm

spacing and cooling rate, two types of molds were employed.

One was a chill mold fabricated from mild steel and the

other was an insulating sand mold. Both molds are shown

in Figure 7. The chill mold was of a split cylindrical

design with the mold cavity being tapered over its 8 inch

height from a diameter of 1 3/4 inches at the top to 1 1/2

inches at the base. This taper was used to help reduce pipe

shrinkage. The walls of the mold were 2 inches in thickness.

The cylindrical mold was set on a 2-inch thick steel base

plate.

The insulating mold was formed over an aluminum

pattern such that the cylindrical taper and wall thickness

were identical to those of the chill mold. The silica sand

was mixed with a binder known as Chem-Rez 270. This binder

was obtained from Ashland Chemical and possessed zero

nitrogen and zero water, The binder was mixed with the

silica sand and Ashland's Chem-Rez Catalyst C-2006. Mixing

time for the sand mixture was approximately 15 minutes.

41

Figure 7: Chill and insulating molds -- The insulating mold is on the left, the chill mold on the right and a typical ingot is in the center.

42

After the sand was properly mixed, it was rammed around the

aluminum pattern. The pattern was stripped in 40 minutes.

It was not necessary to bake the mold since it cured in

air after 2 hours. In order to achieve directional solidi

fication, the cylindrical sand mold was placed on the 2-inch

thick chill plate.

5.4 Casting of the Alloys

The nickel alloys were vacuum cast into either the

cylindrical insulating or chill mold as discussed in section

5.3. For monitoring the temperature of the alloy as it

solidifies, a Pt - Pt/10 Rh thermocouple encased in a thin

quartz tube was inserted into the mold prior to pouring the

metal. In both molds, the thermocouple was located 2%

inches from the base chill plate. A Pt - Pt/10 Rh thermo

couple was also placed in the outer mold wall 2% inches

from the base plate. The output from each of these

thermocouples was fed to Heath Multispeed Strip Chart

Recorders. These recorders had multiranges of 10 mv to 10 V;

a range of 20 mv full scale was selected. This range was

carefully calibrated against a Leeds and Northrup potentiom

eter. In most cases, a chart speed of 2 inches per minute

was chosen.

The melting and casting set-up is shown in Figures 8

and 9.

Figure 8: Melting set-up inside vacuum chamber.

44

Figure 9: Control panel for vacuum melting apparatus.

I

45

5.5 Electron Microprobe Analysis

The purpose of the electron microprobe analysis was

to determine the degree of microsegregation in the casting.

Analyses were made in an Applied Research Laboratory Scan-

ing Electron Microprobe Quantometer (SEMQ); it is shown in

Figure 10. This instrument, as the name implies, is a

combination microprobe and scanning electron microscope

(SEM). The size of the electron beam normally used in this

analysis was 1 micron; although the beam size could be

enlarged to 100 microns. In addition, the microprobe has

the ability to analyze either by x-ray energy or wavelength.

Castings were cut into %-inch thick cylindrical

sections; these samples were then ground and polished in

preparation for analysis. The cylindrical sections were

taken immediately below the plane which contained the tip

of the Pt - Pt/10 Rh thermocouple used to monitor the

solidification rate. In the case of those alloys cast in the

insulating mold, in addition to the aforementioned specimen,

a specimen adjacent to the chill base plate was evaluated.

See Figure 11 for further clarification on the location of

the SEMQ specimens.

The SEMQ was used to make a point-by-point

quantitative analysis across and between dendrite arms.

The wavelength dispersion method was employed in these

determinations. In this type of analysis, the microprobe

Figure 10: Scanning Electron Microprobe Quantometer (SEMQ)

47

Thermocouple

Mold

SEMQ Specimen

-< i—

Ingot

Figure 11: Schematic showing location of SEMQ specimen.

48

uses x-ray spectrometers to measure, identify, and count

x-rays based on their wavelength. For a given point, the

x-ray intensity of each element is compared to the intensity

of the x-rays obtained from the pure standards. The

standards employed in this case were those from which the

alloys were prepared. To convert observed x-ray intensity

ratios into true weight percents, it was necessary to apply

three correction factors to the intensity ratios. These

factors were 1) atomic number correction, Z, 2) absorption

correction, A, and 3) fluorescence correction, F. A program

known as the ZAF program was used to convert the intensity

ratios into weight percent. The ZAF program is presented

in Appendix B.

The energy dispersion method was not employed for

quantitative analysis but was used for qualitative analysis.

In the energy dispersion method, instead of using a crystal

to disperse the emitted x-rays, according to wavelength,

and counting each wavelength interval separately, the x-rays

are picked up directly by a counter which converts them to

pulses with an energy distribution proportional to the wave

length of the x-rays. The energy distribution was displayed

on an oscilloscope screen and the elements present at a

given point were visually determined.

Normally a sample larger than 1 inch in diameter

cannot be placed in the SEMQ. However, a special aluminum

49

block was fabricated and located in the SEMQ which allowed

samples as large as 2 inches in diameter to be used. Thus,

cylindrical cross sections from the casting could be put

into the SEMQ and analyzed.

5.6 Metallographic Analysis

All the castings were sectioned and examined

metallographically. Metallography was used to help

characterize microstruetural features, such as dendrite arm

spacings, phases present, grain size, and the presence of

inclusions. The sections selected for metallographic study

were wet ground on silicon carbide abrasive paper, ranging

in fineness from 240 to 600 grit. Fine polishing of the

specimens was accomplished using 1 micron diamond compound,

followed by polishing with 0.25 micron diamond paste.

Two types of etchants were used. Type I etchant was

50 HNC>g-50 acetic acid. This etchant was prepared fresh

daily with colorless nitric acid to avoid staining the

specimen. The samples were swabbed with the etchant 5-20

seconds, then immediately rinsed with distilled water.

Following the distilled water rinse, the samples were washed

with alcohol. The nitric-acetic acid mixture was used to

delineate the dendrite structure.

To observe grain boundaries, the Type II etchant was

employed. The composition of the etchant was 70 ml

HC1-10 ml H2O2 (307o) - 2 drops of HF as an activator. The

50

samples were immersed in the etchant for 10-90 seconds,

then immediately rinsed with distilled water followed by

final alcohol rinse.

After etching, the samples were analyzed and

photomicrographs were taken using a Reichert He F2

metallograph.

5.7 Differential Thermal Analysis

Differential thermal analysis (DTA) was employed to

help define the nonequilibrium liquidus and solidus tempera

tures for the various alloy composite. In this technique,

the sample temperature is continuously compared with a

reference material temperature. The DTA was performed by

Special Metals Corporation. They used a Dupont 990 Thermal

Analyzer. The DTA method employs a resistance furnace with

a Pt - Pt/13 Rh thermocouple adjacent to the heating element

feeding a signal to the program controller which in turn

regulates power to the furnace. Heating and cooling rates

of 10°C/min were used.

The sample size analyzed was approximately 170

milligrams. The specimen to be evaluated was placed in a

AI2O3 crucible. The apparatus was so constructed that the

bottom of the A^O^ crucible was in intimate contact with

a Pt - Pt/13 Rh thermocouple. The reference side was an

AI2O3 crucible containing platinum. The system was filled

51

with argon and an argon flow rate of 150 cc/min was

established. Argon provided a desirable heat transfer

medium and presented oxidation of the specimen.

For each sample analyzed, a thermogram was made.

On the thermogram, temperature of the sample is recorded

on the ordinate as a function of heat absorbed or

released along the abscissa. In this manner, any phase

transformation which occurred was recorded on the graph.

CHAPTER 6

RESULTS AND DISCUSSION

To analyze the solidification behavior of nickel

alloys containing titanium and aluminum, to establish the

effect solidification rate has on microstructure, and to

determine the effect of alloy melt-crucible interactions

on inclusion formation, eighteen alloys were prepared

and cast. The variables employed for each of these alloys

is given in Table 3. These variables include the type

of crucible in which the alloy was melted, the type of

mold in which the alloy was cast, and the local

solidification cooling rate.

After the alloys were cast, the compositions of the

alloys were verified by spectrographic analysis. The

composition of each cast alloy as determined by

spectrographic analysis is given in Table 4.

In order to determine the degree of microsegregation,

thirteen of the eighteen cast nickel alloys were subjected

to electron microprobe analysis. The microsegregation data

for each alloy analyzed are given in Tables C-l through C-20

of Appendix C. This information is provided for those who

may need this data for other investigations, such as

macrosegregation studies of these alloys.

52

1

2

3

4

5

6

7

8

9

10

11

12

13

TABLE 3

CASTING VARIABLES

Nominal Composition Melting Crucible Casting Mold Cooling Rate

Ni A1 Ti MgO ^•'*2^3 Chill Insulating °C/sec

93 5 2 X X 2.7

93 5 2 X X 0.55

90 10 X X 5.5

90 10 X X 2.7

90 10 X X 0.56

89.5 8.5 2 X X 2.7

89 .5 8.5 2 X X 0.46

90 5 5 X X 3.1

90 5 5 X X 0.55

90 2 8 X X 2.7

90 2 8 X X 0.56

90 10 X X 3.0

90 10 X X 0.55

TABLE 3 (Continued)

Sa™Ple Nominal Composition Melting Crucible Casting Mold Cooling Rata

Ni A1 Ti MgO Al^O^ Chill Insulating °C/sec

14 90 5 5 X X 2.8

15 89 9 2 X X 2.8

16 93 5 2 X X 2.8

17 100 X X 2.7

18 100 X X -

TABLE 4

SPECTROGRAPHIC ANALYSIS OF CAST ALLOYS

Sample Composition No

Ni w/o

A1 W/o

™Ti w/o Fe

w/o Si

w/o C

w/o Mn

w/o Mg ppm

0 ppm

N ppm

Sn ppm

1 92.8 5.2 2.0 .015 - .006 - 6 2 <20

2 92.4 5.6 1.9 .05 .04 .072 .02 59 8 1 <20

3 90.1 9.9 - .03 .02 .007 .01 50 13 - <20

4 90.0 10.0 - .02 .02 .006 .01 50 12 3 <20

5 89.6 9.4 0.6 - .01 .0048 <.01 29 14 2 <20

6 89.0 9.5 1.5 .03 .03 .006 .01 50 12 2 <20

7 89.1 9.3 1.6 .04 .03 .004 <.01 44 10 2 <20

8 89. 7 5.2 5.1 .02 .02 .006 .01 50 12 4 <20

9 89.9 5.1 5.0 .05 .02 .003 .01 48 7 8 <20

10 90.55 2.45 7.0 <.01 .01 .003 .01 60 30 5 <20

11 90.1 2.4 7.5 <.01 <.01 .004 <.01 58 19 6 <20

12 90.0 0.01 9.95 - <.01 .0023 <.01 58 91 8 <20

TABLE 4 (Continued)

Sample Composition No Ni

w/o A1

w/o Ti

W/o

a) o

Si w/o

c W ' / o

Mn w/o

Mg ppm

0 ppm

N ppm

Sn ppm

13 90.0 0.01 9.90 - <.01 .0030 <.01 59 53 4 <20

14 90.1 5.2 4.6 .04 .01 .0028 .01 49 20 3 <20

15 89.0 9.2 1.75 - .01 .005 - 42 18 1 <20

16 93.4 4.6 1.9 - .02 .0042 - 49 13 2 <20

17 99.93 0.1 0.5 <.01 <.01 .0109 <.01 26 10 2 <20

18 100.0 _ _

57

6.1 Nickel-Aluminum and Nickel-Titanium Phase Diagrams

Before discussing the microstrueture of these nickel

alloys, a review of the Ni-Al and Ni-Ti binary systems is

presented. The Ni-Al binary as given by Hansen (27,p.119) is

given in Figure 12. The Ni-Ti binary system according to

Hansen (27,p.1050) has two versions and they are both pre

sented in Figures 13 and 14. Since the alloys of interest

for this investigation contained approximately 90w/o nickel,

only the nickel rich end of the diagram will be considered.

The nickel rich end of the Ni-Al diagram contains

the incongruent melting compound, Ni^Al, and the eutectic

reactions, liq. t Y+Y ' , at 12.6w/o A1 and at a temperature of

1385°C. The y' phase is Ni^Al and the y phase is nickel

solid solution. From the phase diagram, it can be seen that

the Y'-Y eutectic is degenerate. This means that the y-y'

composition lies within the y' phase field at lower

temperatures. Therefore, at approximately 1250°C, the y-y'

eutectic converts to y'. The major phase of this eutectic

is y' and the y1 appears in the form of large particles

separated by thin lamellae of y phase.

The amount of aluminum which can be dissolved in

the nickel-aluminum solid solution is a maximum of llw/o A1

at the eutectic temperature and decreases to 4w/o A1 at room

temperature. Another point of interest is that alloy

compositions between pure nickel and the y-y' eutectic have

a very small equilibrium solidification range (<15°C).

58

10 20 30 40 J I L

WEIGHT PER CENT NICKEL 50 60 70 80 J I L.

1638°

\ 1385° 79(891 183.3)

( N i A l )

2.7 5.7) 72.51 177

(85.1)1 1(87.8) 89.5 (94.8)

300

200

(00

1 1 i I / ! i i i

!

/ I /S>

1! 1 k 1

! i-1 1 1 1

.221 * I 1 i

0 •Al

10 20 30 40 50 60 ATOMIC PER CENT NICKEL

70 80 90 100 Ni

Figure 12: Nickel-aluminum binary system.

59

WEIGHT PER CENT NICKEL

1700

o COOLING * HEATING

THERMAL ARREST, REF. 20

1600

1500

MOO 1380°

1200

57.0\ ..V .(61.9lt--.6lT

/ (-66.2)

(Ni) 1100

1000 (36.51

900 53.1

800

700 0 10 20 10

ATOMIC PER CENT NICKEl 50 30 60 80 70 100

Figure 13: Nickel-titanium binary system: Version I.

60

WEIGHT PER CENT NICKEL 10 20 30 40 SO 60 70 80 90

1800

a ONE PHASel-p... x TWO PHASEJ-• REF. 12,13 * REF. 15

1720®

400 1700

'V MAGNETIC ^.^TRANSFORMATION

"• MARIAN, REE6 200 1600

1500

AT.-% Ti

1400

1300 1287' aoa 85

(8,3.81 \ .1240°

1200

66.8 (89) 1110°

60165)

1015° (37.5)

1000

24.5 (28.5)

11(13)

900

800

90.4 (92)

4(5)

700

600

500 100 90 20 40

ATOMIC PER CENT NICXEL 50 60 80

Figure 14: Nickel-titanium binary system: Version II.

61

Consider the room temperature equilibrium micro-

structure of nickel-aluminum alloys containing between

12.6w/o A1 (eutectic composition) and pure nickel. Alloys

from 12.5 to 12.2w/o A1 will consist of Y' surrounding the

Y phase since the Y~Y' eutectic is degenerate in this

r e g i o n . A l l o y s f r o m 1 2 . 2 t o l l w / o A l w i l l c o n s i s t o f Y

primary phase and Y~Y' eutectic. Alloys between 11 and

4w/o Al will consist of Y primary phase containing Y'

precipitate. The y' precipitate has a cubic or globular

form, as shown in Figure 15. Alloys between 4 and 0w/o Al

w i l l c o n s i s t o f o n l y t h e s i n g l e p h a s e , y .

The nickel-rich end of the Ni-Ti binary system

contains the compound, Ni^Ti (known as n), which melts

congruently at 1380°C, and the eutectic reaction,

liq.^n+y. The eutectic composition and reaction temperature

are still unresolved. Information obtained from this

investigation will help fix the eutectic composition; this

will be presented later. For now, it should be noted that

the eutectic is not degenerate, as is the Y-Y' eutectic in

the Ni-Al binary system. In addition, the n-phase precipi

tates in the Y phase in an acicular or Widmanstatten

pattern. This is illustrated in Figure 16. The

solubility of titanium in nickel decreases from approxi

mately 12w/o Ti at the eutectic temperature to 9w/o at

room temperature.

62

Figure 15: The y ' precipitate in the y matrix -Magnification 10.000X.

Figure 16: The n phase precipitate in the y matrix Magnification 10,000X.

64

The room temperature equilibrium microstructure of

Ni-Ti alloys between the eutectic composition and 11.6w/o

A1 will consist of primary y phase and the n-y eutectic,

which is Widmanstatten in appearance. Alloys between

88.4 and 91 w/o will consist of primary y and, within the y

phase, the acicular n precipitate. Nickel-titanium alloys

less than 9w/o Ti will be comprised of the single phase, y.

6.2 The Nickel-Titanium-Aluminum System

The inter-relationships between the primary nickel

solid solution (y) and the two neighboring phases, n(Ni^Ti)

and y'CNi^Al), in the ternary system have been studied by

Taylor and Floyd (39). These investigations established

isothermal sections for the Ni-Ti-Al system from 1150°G to

750°C. These isothermal sections are shown in Figures 17

and 18. From these figures, it is observed that the y

phase field remains roughly triangular in shape, shrinking

towards the nickel corner as the temperature falls. The

ternary system also contains a ternary eutectic which

results in a y+y'+n three-phase region. The y apex of

this three-phase field closely approaches the nickel-

titanium side of the composition triangle. The two-phase

field of y+n takes the form of a narrow wedge with its apex

at Ni^Ti. The Y+Y' two phase field is extensive, reaching

at 750°C, almost to the nickel-titanium side of the system

at the Y end, and to about 14w/o titanium at the Y' end.

1000 *c

y r -n®

100

NICKEL, ATOMIC PER CENT.

IISO *c

<> 10

' V-O-O m » NICKEL. ATOMIC PER CENT.

—'Pi 100

Figure 17: Isothermal sections for the nickel-titanium-aluminum phase diagrams at 1150°C and 1000°C.

ON Ui

ISO

NICKEL ATOMIC PER CENT.

20 % C l O

O-O 100

MCKEU ATOMIC PER CENT.

Figure 18: Isothermal sections for the riickel-titanium-aluminum phase diagram at 850°C and 750°C.

<T>

67

The Y' phase, when it occurs in a matrix of y,

can be readily distinguished microscopically from the n

phase by its cubic or globular form, which is in marked

contrast to the acicular form of the Ni^Ti (fi) phase. This

difference in structural appearance was discussed in section

6.1. It is also noted that the y' phase, in contrast to the

n phase, exists over a considerable range of composition.

From Figures 17 and 13 it is seen that the range of

solubility of aluminum in the ternary y' phase is much the

same as in the binary nickel-aluminum system, but according

to Taylor and Floyd (39) , it is possible to dissolve

titanium until three out of every five aluminum atoms are

replaced by titanium.

Although Taylor and Floyd determined isothermal

sections for the nickel-titanium-aluminum ternary, they did

not investigate the liquidus surface. The liquidus surface

is needed to aid in the understanding of the solidification

behavior of these alloys. The proposed nickel-rich corner

of the liquidus surface was constructed using 1) the data

from the above isothermal sections of Taylor and Floyd,

2) data obtained in this study from the differential

thermal analysis of selected ternary alloys and 3) micro-

segregation data as discussed in section 6.4. Table 5

gives the composition of the alloys analyzed and the trans

formations obtained by DTA for these alloys.

68

TABLE 5

DIFFERENTIAL THERMAL ANALYSIS DATA

Composition,

Ni A1

w/o

Ti Y / Y 1

Boundary

Transformations,

Equilibrium Solidus

°C

Liquidus

93 5 2 1381 1413 1420

90 10 1377 1386 1412

89.5 8.5 2 1375 1395 1400

90 5 5 1355 1392 1399

90 2 8 1343 1380 1384

1

69

The segregated liquidus surface for the nickel-

rich corner of the Ni-Ti-Al system is presented in Figure

19. From Figure 19, a vertical portion of the Ni-Ti-Al

system at a constant composition of 90w/o Ni was con

structed. The vertical section in Figure 20 proposes a

trough on the liquidus surface. This trough originates at

the ternary eutectic. The equilibrium solidification

range for ternary alloys containing 90w/o Ni is less than

10°C. However, the nonequilibrium solidification range

for these alloys is approximately 40°C.

The ternary liquidus surface can be used to discuss

the microstruetures obtained in the cast Ni-Ti-Al alloys.

Although as-cast alloys are not equilibrium structures,

equilibrium diagrams can be used as an aid in determining

the nonequilibrium phases present. For example, the

equilibrium partition ratio, k, obtained from an

equilibrium diagram can be used in the Scheil equation

(see section 2.5) to obtain the amount of nonequilibrium

second phase present in the alloy.

6.3 Microstructure of the As-Cast Alloys

Six different alloy compositions, as well as pure

nickel were melted and cast. These six compositions are

plotted on the liquidus surface of the nickel-titanium-

aluminum given in Figure 21. The microstructure of each of

the six as-cast alloys will be individually discussed below;

70

20

131J 1325

1350'

70 80 90 100

w/o Ni

Figure 19: Liquidus surface of the nickel-rich end of the nickel-titanium-aluminum phase diagram --Isothermal lines are shown from 1425 to 1325°C.

71

Liquid

1400-

Y+L

° 1300--

& 1200

1100-'

90Ni 8A1

90Ni 6A1

90Ni 4A1

90Ni 2A1

90Ni lOTi

2Ti 4Ti 6Ti 8Ti

Composition, w/o

Figure 20: Vertical portion of t|je nickel-titanium-aluminum system at 90 /o nickel.

72

LOO

/o Ni

Figure 21: Alloy composition investigated plotted on the liquidus surface of the nickel rich end of the nickel-titanium-aluminum phase diagram --Dots represent alloy composition.

73

however, in general, it can be said that the ternary alloys

consist of y' precipitate in a y dendritic matrix and

Y-y' mixture (in the literature known as the binary

eutectic) in the interdendritic region.

Dendritic growth in these alloys gives rise to

microsegregation that greatly affects the formation and

distribution of a secondary phase. For example, consider

the binary alloy, 90w/o Ni-10w/o Al. The equilibrium phase

diagram indicates that the room temperature microstrueture

of this alloy should consist of y primary phase containing

y' precipitate. However, metallographic and electron

microprobe analysis of this alloy indicates that the

microstructure consists of y' precipitate in cored y den

drites, with additional y' occupying the interdendritic

regions. Thus, as a result of nonequilibrium solidification

a fraction of the liquid reaches the eutectic. As

previously discussed, this eutectic is degenerate; therefore,

the y-y' eutectic converts to the y' phase as a result of

solid state diffusion. Figure 22 shows the nonequilibrium

microstrueture of 90w/o Ni-10w/o Al.

The Scheil equation, Eq. 4 in section 2.5, applied

to the 90w/o Ni-10w/o Al alloy indicates that as a result

of nonequilibrium solidification a eutectic mixture

amounting to 23.6w/o should be present. No solid diffusion

74

Figure 22: Microstructure of 90w/o Ni-10w/o A1 -- The dark phase is cored y dendrites; the white phase is y'. Magnification is 74X.

75

was considered in obtaining the value of 23.6w/o, In order

to semi-quantitatively determine the percentage of

degenerate y-y' eutectic mixture present, the point-count

technique (40) was used. Using this method, it was

determined that 22.1w/o of the alloy consisted of the

degenerate eutectic mixture. This is slightly less than

predicted by the classical Scheil equation.

As previously discussed, the Scheil equation

assumes no solid diffusion. The Scheil equation has been

modified by Flemings (14) in order to account for solid

diffusion; the expressions obtained were:

fs k-1 Cs " *Co C1- Tl4s> <7>

a-fL) k-1

CL " Co (1" 1+ak ) (8)

4DS t f where ' a = -— (9)

d2

The term Dc, in cm /sec, is the diffusion coefficient of u

the solute in the solid, t^, in seconds, is the local

solidification time, and d, in cm, is the dendrite arm

spacing. All the other terms in Equations 7 and 8 were

previously discussed in section 2.5. For this 90w/o Ni-

10w/o A1 alloy, Da, from the data of Sherby and Simnad (41) D -10 2 was approximately 9 x 10 cm /sec, t^ was 6 seconds, and

d was 14 x 10"4 cm; this results in an a of 0.012 and an

76

ctk value of 0.01. Applying the dimensionless a parameter

in Equation 8, the amount of second phase was calculated to

be 22.8w/o. This more closely agrees with the value of

22.1w/o obtained by the point-count technique. For ak<<l,

microsegregation approaches the maximum predicted by the

classical nonequilibrium Scheil equation; for ak<<l the

composition of the primary solid phase approaches uniformity.

Therefore, for this alloy solid state diffusion did not

have a significant effect.

The microstrueture of the 90w/o Ni-10w/o Ti alloy

is shown in Figures 23 and 24. Figure 23 shows a structure

of y dendrites and N-Y eutectic in the interdendritic

regions. Figure 24 shows a magnified view of the needle

like n-y eutectic. As with the 90w/o Ni-10w/o A1 alloy,

the eutectic mixture in the 90w/o Ni-10w/o Ti alloy is a

manifestation of nonequilibrium solidification. The

classical Scheil equation predicts the amount of eutectic

to be 29.7w/o. The actual amount of eutectic as determined

by the point-count technique was 27.1w/o. The amount of

eutectic predicted by the modified Scheil equation was W _ Q 9

28.9 /o. This was using a value of 1 x 10 cur/sec for

Ds, 11 seconds for t£, and 18 x 10"^ cm for d. The

effect of solid diffusion is small, as in the Ni-Al system.

Thus the calculated value of 28.9w/o obtained using the

modified Scheil was in close agreement with the value of

77

Figure 23: Microstructure of 90w/o Ni-10w/o Ti -- The white phase is cored y dendrites, the dark phase is n-y eutectic. Magnification 74X.

78

JILL '

Figure 24: Microstructure of 90 /o Ni-10 /o Ti showing n-y eutectic between y dendrite arms --Magnification 900X.

79

27.1w/o obtained from the point-count technique. In

addition, there is very little n precipitate in the y

dendrites, since an alloy of this composition crosses the

solvus line at near room temperature where solid diffusion

is very slow.

In ternary alloys, as a result of nonequilibrium

soli d i f i c a t i o n , a fr a c t i o n o f t h e l i q u i d r e a c h e s t h e y / y 1

boundary. The y/y' boundary is line ae in Figure 19.

Solidification is complete when all the liquid is exhausted

or when the liquid reaches the ternary eutectic at point e

on Figure 19. Metallographic, electron microprobe, and

differential thermal analysis techniques indicated no

presence of a ternary eutectic in the alloys investigated.

Presumably, in these alloys, all the liquid is exhausted

along the y/y' boundary (line ae in Figure 19) before the

ternary eutectic is reached. All the ternary alloys showed

varying amounts of y-y' in the interdendritic regions.

Using the point-count technique, the approximate amount of

y-y' mixture in each of the alloys was determined. The

resulting data are presented in Table 6. The cooling rate

made only a slight difference in the amount of second phase.

The reason why microsegregation is so nearly constant over

wide cooling rate ranges is that the coarseness of the

dendrite structure, as measured by the dendrite arm spacing,

80

TABLE 6

PERCENTAGE OF y-y1 EUTECTIC MIXTURE IN THE ALLOYS

Mold Composition, w/o Chill Insulating

Ni A1 Ti 7o Second Phase % Second Phase

90 10 22.1 20.6

89.5 8.5 2 70.0 69.1

90 5 5 11.5 11.2 i

90 2 8 39.9 39.6

90 10 27.1 26.8

93 5 2 ~2.0 ~2.0

81

varies with the cooling rate. This variation is such

that the extent of diffusion after solidification is

nearly constant.

Investigation of the 89.5w/o Ni-8.5w/o Al-2w/o Ti

alloy showed that the y in the y-y' mixture which formed as

solidification proceeded along the y/y' boundary (line ae

in Figure 19), degenerated to y'. This is illustrated in

the photomicrograph of this alloy in Figure 25. Since the

composition of this alloy lies very close to the maximum

solubility of aluminum and titanium in nickel, the large

proportion of second phase in this alloy, approximately

70%, is to be expected.

The microstruetures of ternary alloys

90w/o Ni-5w/o Al-5w/o Ti and 90w/o Ni-2w/o Al-8w/o Ti are

shown in Figures 26 and 27, respectively. For these

alloys, not all the y-y' in the interdendritic regions

degenerated to y*. This lack of degeneration is shown in

Figure 28, in which y' phase appears in the form of

large particles separated by lamellae of y phase. The

composition of alloy 90w/o Ni-2w/o Al-8w/o Ti lies closer

to the y/y' boundary than the 90w/o Ni-5w/o Al-5w/o Ti

alloy. Therefore, assuming that the equilibrium partition

ratio, k, for the two alloys is equal, and vertical section

82

Figure 25: Micro structure of 89.5 /o Ni-8.,5 /o Al-2 /o Ti -- The dark phase is cored primary y dendrites, the white phase is y'. Magnification 74X.

Figure 26: Microstructure of 90w/o Ni-5w/o Al-5w/o Ti The light phase is a y-y1 mixture. Magnification 74X.

84

Figure 27: Microstructure of 90w/o Ni-2w/o Al-8 /o Ti--The dark phase is cored Y dendrites; the white phase is a Y~Y' mixture. Magnification 74X.

35

tv.-- r- -

Figure 28: Microstrueture of y-y1 mixture between dendrite arms -- The dark phase is y, the lighter phase is y'. Magnification 5000X.

86

in Figure 20 indicates they would be similar, the

90w/o Ni-2w/o Al-8w/o Ti should contain a greater amount

of nonequilibrium secondary phase. This was found to be

the case, as the point-count technique indicated the

90w/o Ni-2w/o Al-8w/o Ti and the 90w/o Ni-5w/o Al-5w/o Ti

alloys contained approximately 39 and llw/o, respectively.

The 93w/o Ni-5w/o Al-2w/o Ti alloy contained

approximately 2w/o secondary phase. This alloy is much

further from the y/y' boundary than the other ternary

alloys. In addition, from the data in the phase diagrams,

it appears that the liquidus and solidus lines are slightly

steeper for this alloy than the others; this results in a

larger equilibrium partition ratio (k approaches 1). Thus

a smaller amount of secondary phase is present in this

alloy. The microstrueture of this alloy is shown in

Figure 29. Note that no y 1 precipitated in the cored y

dendrites on cooling. An alloy of this composition crosses

the solvus line near room temperature where the rate of

diffusion is slow.

To summarize section 6.3, it can be stated that

as a manifestation of nonequilibrium solidification, all

the cast alloys investigated showed cored y dendrites with

either a secondary phase of y' or a y-y1 mixture occupying

the interdendritic regions. For the binary alloy,

90w/o Ni-10w/o Al, the interdendritic regions contained the

Figure 29: Microstructure of 93w/o-5w/o Al-2w/o Ti The structure consists mostly of cored dendrites. Magnification 74X.

88

degenerate y - y 1 eutectic; for the 90w/o Ni-10w/o Ti alloy,

n-y eutectic was present between the dendrite arms. The

actual percentage of eutectic mixture present in these

binary alloys agreed closely with the values calculated

by the Scheil equation. Modifying the Scheil equation to

account for solid diffusion improved the agreement between

predicted and measured values, but also showed that solid

diffusion did not have a significant effect on the amount

of eutectic mixture in these binary alloys.

The y - y ' mixture in the ternary alloys resulted

when some of the liquid reached the y/y' boundary in the

Ni- T i - A l t e r n a r y s y s t e m . A f t e r t h e l i q u i d r e a c h e d t h e y / y 1

boundary, y-y' formed between the cored dendrite arms as

the liquid moved along the boundary. The y-y' degenerated

to y' in the 89.5w/o Ni-8.5w/o Al-2w/o Ti alloy but did not

degenerate in the other ternary alloys with lower Al/Ti

ratios. Although the nickel-rich end of the Ni-Ti-Al

ternary system contains a ternary eutectic close to the

Ni-Ti boundary, none of the ternary alloys exhibited this

eutectic. Thus, all the liquid was exhausted along the

y/y1 boundary before the ternary eutectic was reached.

The percentage of y-y' mixture in the ternary alloys varied.

The amount of this mixture is a function of the equilibrium

partition ratio, k, and the proximity of the alloy composi

tion to the y/y' boundary.

89

6.4 Dendritic Microsegregation

Microsegregation is the nonuniform distribution of

alloying elements, the period of the nonuniformity being

on the scale of the dendrite arm spacings. In isomorphous

alloys, microsegregation occurs as local minimums and

maximums in concentration. In multiphase alloys, in

addition to coring, microsegregation occurs either as

formation of secondary phases where none is predicted by

the equilibrium diagrams, or microsegregation occurs as

formation of more than the equilibrium amount of secondary

phases where some secondary phases are predicted.

For these Ni-Ti-Al alloys, the variation of the

alloying elements concentrations across the secondary

dendritic arms and the interdendritic regions were

determined using electron microprobe techniques. As

previously mentioned, the microsegregation data for the

alloys analyzed is given in Appendix C. The data in

Appendix C is summarized in Table 7 which gives the

segregation ratio (SR) of titanium and aluminum in the

secondary dendrite arms, between the arms, and the overall

SR, that is, the ratio of the overall maximum to the overall

minimum concentration. If the interdendritic secondary

phase is a binary or ternary eutectic, the composition of

the eutectic will be constant. However, in these nickel-

base alloys, the interdendritic regions have concentration

TABLE 7

SEGREGATION RATIO DATA FOR Ni-Al-Ti ALLOYS

Nominal Alloy Composition Ni A1 Ti

Cooling Rate

°C/sec

SR Arms

Ti A1

SR Between Ti

Arms A1

SR Overall

. Ti A1

93 5 2 2.7 1.61 1.04 1.57 1.03 2.70 1.06

93 5 2 0.55 1.19 1.04 1.40 1.03 1.60 1.05

90 10 5.5 1.40 1.40

90 10 2.7 1.25 1.25

90 10 0.46 1.21 1.21

89.5 8.5 2 2.7 1.68 1.20 1.10 1.02 1.80 1.10

89.5 8.5 2 0.46 1.50 1.10 1.10 1.02 1.60 1.10

90 5 5 3.1 1.68 1.10 1.35 1.20 2.24 1.20

90 5 5 0.55 1.51 1.10 1.20 1.03 2.10 1.10

90 2 8 2.7 1.09 1.08 1.19 1.15 1.43 1.24

90 2 8 0.56 1.07 1.02 1.18 1.14 1.37 1.17

90 10 3.0 1.33 1.33

90 10 0.55 1.22 1.22

91

variation (segregational effects) as a result of the

liquid solidifying as it moves along the y/y' boundary

(line ae in Figure 19). Thus the segregation ratio in

the region between the arms is reported. Examination of

the data in Table 7 indicates the following:

(1) Aluminum shows significant segregation across

the dendrite arms of both chill cast and

slow cooled binary Ni-Al alloys.

(2) The SR of aluminum in the dendrite arms

decreases as the composition of the alloys

moves from the 90w/o Ni-10w/o A1 binary

alloy across the ternary at a constant

composition of 90w/o Ni to the

90w/o Ni-2w/o Al-8w/o Ti ternary alloy.

(3) Titanium shows significant segregation in both

the binary and ternary alloys.

(4) The SR of titanium in the dendrite arms

first decreases as the composition of the

alloys moves from the 90w/o Ni~10w/o Ti binary

alloy across the ternary at a constant

composition of 90w/o Ni to the

90W/o Ni-2w/o Al-8W/o Ti ternary alloy and

then the SR of titanium increases again as

the ternary alloy composition moves toward

the 90w/o Ni-10w/o A1 binary.

92

(5) The segregation of both titanium and aluminum

is slightly greater in the chill cast than in

the slow cooled alloys.

The SR of aluminum in the dendrite arms is 1.4 for

the chill cast binary alloy, 90W/o Ni-10W/o Al, moving into

the ternary system from the binary by adding titanium

while maintaining the composition of nickel at approximately

90w/o causes the SR for aluminum in the arms to decrease to

1.08 for the chill cast 90w/o Ni-2w/o Al-8w/o Ti alloy.

On the other hand, the SR of titanium in the dendrite arms

decreases from 1.33 in the binary to 1.09 for the chill cast

90w/o -2w/o Al-8w/o alloy and then increases again to 1.68.

Presumably, this occurs as a result of the ternary eutectic.

As shown in Figure 20 (in section 6.2), a trough is proposed

in the 90w/o Ni vertical section. This trough has a temper

ature minimum at an approximate composition of 90w/o Ni-

0.5w/o Al-9.5w/o Ti. As alloys approach this composition

along the 90w/o Ni vertical section, the equilibrium parti

tion ratio approaches 1. According to the classical Scheil

equation, as k approaches 1, the SR approaches 1 for both

aluminum and titanium. Moving away from the 90w/o Ni-

10w/o Ti binary composition, into the ternary system along

a path such that the nickel composition remains constant,

will first cause the SR of the alloying elements in the

dendrite arms to decrease and then increase again as the

93

composition of the alloy moves past the concave

deflection in the liquidus.

The modified Scheil equation, Equation 8,

showed that microsegregation will depend on the a-factor,

Equation 9. According to the modified Scheil equation,

the case segregation should approach equilibrium for large

values of a. The smaller the value of a, the greater the

segregation ratio up to the limiting case of negligible

diffusion in the solid phases. As discussed in section

6.7, the a factor changes slowly because the dendrite arm

spacing is proportional to the square root of solidification

time, Thus, the extent of diffusion occurring after

solidification is nearly constant. Therefore, as

predicted, for a given Ni-Ti-Al alloy, only a slight

difference in SR was experienced between the chilled and

slow cooled alloys.

The aluminum concentration was slightly erratic

near the dendrite arm/interdendrtic region. The micro-

probe specimens were lightly etched prior to analysis so

the microstrueture could be clearly delineated. This

caused the dendrite arms to be slightly raised above the

interdendritic region. Since the x-ray of a light element

such as aluminum is easily absorbed, surface roughness

slightly effects the absorption coefficient which is used

94

in the ZAF program to calculate concentrations. This

results in slightly erratic results near the boundary.

Titanium is a heavier element and this edge effect was not

experienced.

In order to verify that etching did not

significantly change the concentrations obtained by the

microprobe, etched specimens had their dendrite arms

marked by microhardness indentations and these specimens

were repolished. After repolishing, the microhardness

indentations were still present such that the location of

the arms were known. Microprobe analysis across these

arms showed no significant difference in the alloy

concentrations between the unetched and etched specimens

of identical compositions.

Recently the microsegregation in cast Inconel 713

(Ni-125Cr-4.2No-2.0Cb-6.1Al-0.8Ti) was investigated by

Bhambri, Kattamis, and Morral (42). They reported that

aluminum had a dendrite arm SR of 1.046 and titanium an

.SR of 1.080. They also found that the SR for chromium,

niobium and molybdenum were near unity. Furthermore,

these investigators reported that the aluminum content of

the secondary arm decreased from the arm center to the

edge. They did not indicate if the alloy was etched or

unetched prior to microprobe analysis nor did they

present any of their data. It is possible, therefore,

95

that if they used etched specimens, they had a surface

effect, as discussed above.

In section 6.1, two versions of the Ni-Ti binary

system were presented. One version gave the eutectic

reaction, 5,?n+Y, at the composition of 83.8w/o Ni, the

other diagram gave the composition of the eutectic as

87.4w/o Ni. Microprobe analysis of the 90w/o Ni-10w/o Ti

alloy indicated the interdendritic region contained a

eutectic with a composition of approximately 12.6w/o Ti;

therefore, this investigation indicates the n-y' eutectic

composition is 87.4w/o Ni-12.1w/o Ti.

6.5 Macrosegregation

A formal study of macrosegregation was not

conducted in this investigation; however, in the course

of the microsegregation analysis, some information regarding

macrosegregation was obtained. First, spectrographic

analysis showed no difference in the metal solute or oxygen

concentration between the center and edge of the casting.

Secondly, a point-by-point analysis using the electron

microprobe was made using a beam diameter of 100 microns to

achieve a macroscopic analysis. No quantitative results

were obtained but qualitative results were gathered and

are presented in Figure 30. Only small variations of

aluminum and titanium were noticed during the traverse from

casting edge to center edge.

Figure 30: Macroscopic Scan from the casting edge to casting center for nickel, aluminum and titanium.

97

6.6 Variation of Secondary Arm Spacing with Cooling Rate

The variation of secondary dendrite arm spacings

with local cooling rate for the Ni-Ti-Al alloys investigated

is illustrated in Figure 31. For the various alloys, the

arm spacings remained fairly constant for a given cooling

rate. This is attributed to the fact that the liquidus

temperatures of the alloys are in close agreement.

Therefore, the dependence of secondary arm spacing, d, on

local cooling rate, GR, may be expressed for these alloys

by the single relationship:

d = A (GR)~n =27.7 (GR)"0'38 (10)

The A and n constants were obtained by least squares

analysis.

Differential thermal analysis showed that all of

these alloys have a nonequilibrium solidification range of

between 35 and 40°C. Knowing the nonequilibrium

temperature range (AT) and the local cooling rate (GR),

the local solidification time, t£, was found using the

expression:

tf = 5E (11)

Knowing this information, the variation of arm spacing

was related to local solidification time, t^, by expression:

100 -•

co a o u o

6 C

•<-i

id

W) 0 •H o a w S V-t aj

-0.38 d = 2 7. 7 (GR)

1 10

local cooling rate, GR, °C/sec

100

Figure 31: Variation of secondary dendrite arm spacing with local cooling rate.

vo CO

99

d = 7.0 t°-37 (12)

This relationship is illustrated in Figure 32.

These expressions compare closely to those

previously obtained by investigators for other alloy

systems. Bower, Brody, and Flemings (43) obtained the

expression, d = 7.5 t^'39, for Al-4.5w/o Cu. Bhambri,

Kattamis, and Morral (42) for the cast alloy, Inconel 718C, A / O

found the expression, d = 6.79 tf ' . Brower and

- 0 ^ 2 Flemings (44) obtained the expression, d = 60(GR) ' ,

for Fe-25w/o Ni alloys. Therefore, the expression obtained

in this study relating the dendrite arm spacing to local

solidification rate agrees closely with results of other

studies on a wide variety of alloys (45), as illustrated in

Figure 33.

6.7 Alloy Melt-Crucible Interactions

Since inclusions may form from oxygen dissolved in

the melt upon solidification, or result from particles of

insoluble refractory removed from the crucible, the

reactions between alloy melt and crucible must be studied

in order to determine how best to minimize the total

oxygen pickup. Although vacuum melting prevents oxidation

from the atmosphere of such elements as titanium and

aluminum in the nickel melt, reactions between these

reactive constituents and the refractory materials

comprising the crucible may be facilitated in a vacuum.

•r-l

o Pu CO

e n cO

0 . 1 1 1 0 1 0 0

local solidification time, t^, sec.

Figure 32: Variation of secondary dendrite arm spacing with local solidification time. H-1

o o

100 ••

10

Sn-13 /c Inconel 713C This investigation (90 /o Ni-Al-Ti) Al-4.5w/o Cu Fe-lCT/o Ni Ti-10 /o Fe Ti-10 /o A1

iio 6.1 10

local solidification time, tf) sec.

Figure 33: Secondary dendrite arm spacing as a fmction of local solidification time for several alloys.

102

If these reactions occur, the beneficial effects of

vacuum melting nickel-base alloys are partly offset.

Therefore, as part of this study, Ni-Ti-Al alloys were

melted in both MgO and crucibles, and the subsequent

castings were analyzed for the types of inclusions present.

The crucibles were also examined to determine the amount

of nickel penetration.

One way in which the crucible could contribute

oxygen to the melt, which could eventually result in

oxygen inclusion formation when the solubility limit is

exceeded, is by the decomposition of the refractory

crucible to give oxygen dissolved in the nickel-base alloy

plus a volatile constituent which is carried off in the

vacuum system. First the decomposition of MgO will be

considered.

A possible decomposition reaction for MgO is:

MgO t Mg(g) + 0 (13)

K = (Pmg)(ao> (14) a mgo

In order to determine the partial pressure of the Mg gas,

it is necessary to know the equilibrium oxygen content in

the melt.

Figure 34 gives the equilibrium relationship between

dissolved oxygen content in the nickel melt and the content

of the various deoxidizing elements which the melt contains.

10°

Ti

.012 atm CO

.07 atm CO

wt%

alloying

addition A1

10"

1.31 x 10 atm CO

10

wt % 0

Figure 34: Equilibrium relationship between dissolved oxygen in the nickel melt and the content of various deoxidizing elements.

o LO

104

These relationships were obtained from the thermodynamic

data discussed in Chapter 3 and given in Appendix A. The

carbon-oxygen relationship is shown for three different

partial pressures of CO. The partial pressure of CO is — 6 — o

equal to the vacuum pressure of 1.31 x 10 atm. (1 x 10

torr) at the top of the melt. At this pressure the carbon

becomes a stronger deoxidizer than either titanium or

aluminum. Therefore, the oxygen potential near the top

of the melt will be determined by the carbon content, even

down to very low (<<0.001w/o C) levels.

Analysis of the castings indicate that the alloys

contained approximately 0.0035w/o C, consider the reaction:

C + 0 = C0(g) (15)

K = (16) aoac

Assuming that the activities of 0 and C are equal to their

respective concentrations in weight percent (Henrian

behavior), for an equilibrium concentration of 0.0035w/o C

- fi and a partial pressure of CO of 1.31 x 10 atmosphere,

according to thermodynamic calculations the dissolved

equilibrium oxygen content of the melt is 3.11 x 10~ w/o.

Using this value of dissolved oxygen in Equation 14, a O

partial pressure of Mg of 5 x 10 atm. is obtained. This

is considerably higher than the vacuum system pressure of — fi

1.31 x 10" atm.; therefore, it is seen from these

105

calculations that the MgO crucible will dissociate according

to Equation 13 at the melt-crucible-vacuum interface.

Hence, Mg is constantly pumped from the system with a

consequent input of oxygen into the top of the melt.

In the case of MgO, for every atom of Mg lost to the

vapor, one atom of oxygen goes into the melt.

As a consequence of the metallostatic pressure head, _ £

the pressure head increases from 6.6 x 10 atm. at the

top of the melt to 0.7 atm. at the bottom of the melt

(the height of the melt was normally 3 1/2 inches). As

seen from Figure 34, at a pressure of 0.07 atm.,

aluminum is a stronger deoxidizer than carbon for any

carbon content less than 0.7W/o- For alloys studied in

this investigation, the initial carbon was low, usually no

greater than 0.006w/o, and the aluminum content was fixed

in the melt by alloy additions of from 1 to 10w/o Al.

Thus, for alloys with these compositional ranges,

aluminum will control the soluble oxygen content

immediately away from the top surface toward the bottom

of the melt.

Assuming that the reaction MgO + C -* C0(g) + ®(g)

occurs at the metal/crucible interface, Figure 35 shows

the equilibrium partial pressure of Mg^ in the gas as

a function of w/o C in the melt. This pressure is equal

to the CO pressure and is one-half the total pressure.

-4 10

w/o C

Figure 35: Equilibrium partial pressure of Mg in the gas as a function of w/o C in the melt. i-1 o cr»

107

Figure 36 shows the metallostatic head pressure as a

function of depth below the top surface. From these

figures, it can be seen that a MgO crucible containing a

melt which has a dissolved carbon content of 0.003w/o,

will dissociate via the reaction MgO + C -> + S(g)

to a depth of only 0.0175 inches. Thus, as oxygen

enters the melt near the top from the decomposition of

MgO, it can be inductively stirred away from the top of

the melt to a depth where the dissolved equilibrium

oxygen concentration is controlled by the aluminum content.

When this equilibrium oxygen concentration is exceeded,

AI2O2 inclusions nucleate and form in the melt as

primary inclusions.

One indication that a primary oxide phase was

forming while the nickel alloy melts were held in the

MgO crucible was the observation that the surface of the

melt showed an increasing amount of floating particles

with time. Figure 37 shows the surface of the liquid

alloy in the MgO crucible immediately after melting.

Figure 38 shows the surface of the same alloy melt five

minutes later. Note the increase of the floating particles

on the surface. Electron microprobe analysis revealed

these particles to be A^O^.

1

- 2

3

4 ~-2

depth below top surface, inches

Figure 36: Metallostatic head pressure as a function of depth below the top surface of the melt. o

00

109

Figure 37: Surface of the alloy melt immediately after melting -- Melt composition 90w/o Ni-5w/o Al-5 /o Ti. MgO crucible. Temperature 1500°C.

110

pppopip p

Figure 38: Surface of the alloy melt 5 minutes after melting -- Melt composition 90w/o Ni-5w/o Al-5w/o Ti. MgO Crucible. Temperature 1500°C.

Ill

These AÔg particles tend to agglomerate and

form large clusters. Figure 39 shows such an AI2O3 cluster.

Figures 40 through 42 are characteristic x-ray images of

nickel, aluminum, and oxygen respectively for this cluster.

This verifies that these clusters are AlÔ^. No titanium

was ever observed in such clusters. Braun, Elliott, and

Flemings (46) have found clusters of AÔ^ in steel melts.

AÔg particles rise in an iron or nickel bath because

their density is less than that of the metal and because

of convective stirring. On rising through the melt,

these- particles collide and adhere, with the results

being large interconnected clusters. These clusters

contain a large number of individual inclusions. Also,

during induction melting, the stirring pattern is down

at the center and up along the sides with a region at

mid-radius with low velocity. Therefore, the low density

inclusions will tend to be concentrated at this point

increasing their probability of collision.

For those alloys which were melted in MgO

crucibles, particles of MgO were sometimes found in the

dendrites. This would be expected since as the crucibles

dissociate particles of MgO will fall into the melt. A

photograph of such a particle is shown in Figure 43.

The following figures, Figures 44 through 48, are the

112

Figure 39: Backscattered x-ray image of typical A190~ clusters found in 90w/o Ni-10w/o Al casting--Magnification 2000X.

113

Figure 49: Distribution image of Ni-Ka radiation of A1^0„ clusters in Figure 39 -- Superimposed on tne x-ray image is the relative nickel concentration across the traverse indicated by the upper straight white line. The bottom straight line represents zero nickel concentration. Magnification 2000X.

114

Figure 41: Distribution image of Al-K„ radiation of AI2O0 clusters in Figure 39 -- Superimposed on tne x-ray image is the relative aluminum concentration across the traverse indicated by the upper straight white line. The bottom straight line represents zero aluminum concentration. Magnification 2000X.

115

Figure 42: Distribution image of 0-Ka radiation of AI2O0 clusters in Figure 39 -- Superimposed on tne xpary image is the relative oxygen concentration across the traverse indicated by the straight white line. No base-line concentration for oxygen is given since the concentration of oxygen in the matrix approaches zero. Magnification 2000X.

116

Figure 43: Backscattered x-ray image of MgO particle found in an alloy melted in an MgO crucible -Magnification 2000X.

117

Figure 44: Distribution image of Ni-Ka radiation of particle in Figure 43 -- Dark indicates the

, lack of Ni present. Particles found in 90w/o Ni-Al-Ti alloys melted in MgO crucibles. Magnification 2000X.

118

r;ct/u

Figure 45: Distribution image of Al-K.a radiation of particle in Figure 43 -- Note absence of A1 in center of particle. Magnification 2000X.

119

Figure 46 Distribution image of 0-Ka radiation of particle in Figure 43 -- Magnification 2000X.

120

Jr*

Vk t: 1

- • ' . j ' 10 ••1 v • '•• •i.ynU;

>. r , • :' •" •••» • •' ; •-•••. 'v-i.- 'v.:<;-' • ' . .')* •' .-V ;»v,v.

•• r ••' • / !<•*' v>W • >V*V ..V'

:;'-K'*:ft%r;;::,''f;

J teaarifeflSBI

Figure 47: Distribution image of Mg-Ka radiation of particle in Figure 43 -- Note Mg concentration is greater in the center. Magnification 2000X.

121

Figure 43: Distribution image of Ti-Ka radiation of particle in Figure 43 -- Note particle showed no presence of titanium. Magnification 2000X.

•f

122

related x-ray distribution images of nickel, aluminum,

titanium, oxygen, and magnesium respectively. This sequence

of photographs shows that the MgO reacts with the aluminum

and dissolved oxygen in the melt to form the spinel compound

MgAl20^ around the original MgO particle. As the reaction

continues, MgO diffusion is not fast enough to prevent

further reaction with A1 and 0 to form A1203; careful

examination of the photographs shows that the MgAl20^

layer around the MgO is itself surrounded by a layer of

^2^3• evi-dence of any NiAl20^ formation is observed,

nor was any observed in connection with any other

inclusions studied.

Next, consider the decomposition of the A1203

crucible. Four possible decomposition reactions exist:

A1203 Z 2A1 + 30 (17)

A1203 t Al20(g) + 20 (18)

A1203 t 2A1 0(g) +0 (19)

A1203 t 2Al(g) + 30 (20)

The first reaction is the reverse of the de'oxidization

reaction. If the removal of A1 from the melt by

vaporization is added to it, the result is Equation 15.

As previously discussed, at the top of the nickel alloy

melt a dissolved equilibrium oxygen concentration of

123

3.11 x 10~®w/o is present if the carbon content is 0.003w/o.

For this oxygen concentration, with the aid of the

thermodynamic relationships given in Appendix A, the

corresponding partial pressure of AÔ^g^, AlO^ , and

Al(g) were found to be 1.21 x 10"^ atm., 7.70 x 10 ^ atm., £

and 8.33 x 10~° atm. respectively. The partial pressure

of AÔ^g^ and exceed the vacuum system pressure

of 1.31 x 10 atm., although not to the extent that the

Mg(g) partial pressure exceeded the vacuum system pressure

leading to the decomposition of MgO. Therefore, using the

same line of reasoning as discussed for the MgO

dissociation, oxygen can be pumped into the melt, as a

result of the AI2O3 dissociation to AÔ^ and/or Al^g-j .

The dissolved oxygen can be inductively stirred away from

the top of the melt to a depth where it reacts with the

aluminum in the melt. As in the case of melts in the

MgO crucible, this leads to the formation of AÔg

inclusions which float to the surface of the alloy melt.

Since the partial pressures of AlÔ^^ and Al^ are only

slightly greater than the vacuum pressure, the

dissociation of the AÔg is restricted to the metal

vacuum interface as a result of the effect of the

metallostatic head pressure.

124

Alloys melted in crucibles, contained large

angular particles of AI2O3 in the dendrites. Since these

particles are indicative of those found in the AÔ^

crucible, it can be assumed that these inclusions are

from the crucible itself, rather than as a result of a

reaction between dissolved oxygen and aluminum. A

photograph of such an angular particle, is shown in

Figure 49. Figures 50 through 53 are x-ray distribution

images which verify that the particle is AÔg.

Other evidence that the MgO and AÔg dissociated

during melting of the nickel alloys was obtained from

examination of the crucibles themselves. The Ni-Ti-Al

alloys did not penetrate the MgO or AÔ^ crucibles. Cross

sections of the crucible walls were analyzed by

metallographic and electron microprobe techniques and no

penetration of nickel, titanium or aluminum from the

nickel alloy into the crucible walls was evident. To

further verify that the liquid metal did not penetrate

the crucible walls, fracture surfaces of the crucible

wall cross sections were examined with the scanning

electron microscope (SEM) and in conjunction with the SEM,

the energy dispersive unit of the electron microprobe was

used. These techniques again indicated no metal

penetration. However, both crucibles showed a reduction

in wall thickness at the melt/vacuum interface.

125

Figure 49: Backscattered x-ray image of particle found in an alloy melted in an A1„0„ crucible -- Magnification 1000X.

126

Figure 50: Distribution image of Ni-K radiation of Alr,0q particle in Figure 49 -- Dark areas indicate the lack of Ni present. Magnification 1000X.

127

Figure 51: Distribution image of A1-IC„ radiation of AI9O3 particle in Figure 4y -- Bright area indicates a high concentration of A1. Magnification 1000X.

128

Figure 52: Distribution image of 0-Ka radiation of Al^O^ particle in Figure 49 -- Bright area indicates the presence of oxygen. Magnification 1000X.

129

Figure 53: Distribution image of Ti-K radiation of particle in Figure 49 --

Magnification 1000X.

130

This reduction amomted to approximately 22% of the wall

thickness for the MgO crucible and 157o for the

crucible. Both of these crucibles had been used to prepare

ten nickel alloys; thus these reductions were a result of

the crucible being in contact with the Ni-Al-Ti melts for

a total of 50 minutes. Thus, these data showed the

dissociation occurred near the melt/vacuum boundary, as

predicted by thermodynamic calculations (see Figure 54).

Thus far, only a discussion of primary inclusion

formation has been presented. However, secondary inclusions

do form in these alloys. As discussed in section 4.2,

secondary inclusions result because the solute elements

are rejected to the interdendritic spaces during solidifi

cation. If the oxygen and metal solutes (titanium and

aluminum) being rejected to the interdendritic region

reach sufficient concentration in that region to cause,

supersaturation with respect to a thermodynamically stable

oxide, an interaction will occur and a second phase forms.

The secondary inclusions found in these alloys were small,

ranging from 0.2 to 3 microns. These inclusions are

found between dendrite arms. This li illustrated in

Figure 55 where the secondary inclusions are clearly

identifiable in the microstructure of the 93w/o Ni-

5w/o Al-2w/o Ti alloy.

131

cross section of crucible wall

} melt-vacuum boundary region

Figure 54: Schematic of a crucible wall cross section showing reduction of wall thickness near the melt-vacuum boundary -- Pocked areas of this nature were noted throughout the melt-vacuum boundary region.

132

Figure 55: Primary inclusions in dendrites of 93w/o Ni-5w/o Al-2w/o Ti.

133

Electron microprobe studies showed that the small

secondary inclusions in the Ni-Ti-Al ternary alloys were

AI2O3. TiC>2 inclusions do not form in these ternary alloys

because the deoxidation constant of AI2O3 is much greater

than TiC>2 • To illustrate this point consider the ternary

alloy, 90w/o Ni-5w/o Al-5w/o Ti, which had one of the

highest segregation ratios of titanium between the arms.

In this alloy, the maximum concentration of titanitim

between the arms was 8.40w/o and the maximum concentration

of aluminum was 5.45w/o. Thus for the reactions:

2A1 + 30 = A1203 (21)

aAl?0o K = 2 3 <22>

(aAi)2(a0)3

and

Ti + 20 = Ti02 (23)

aTi00 K = 1— (24)

(aTi)(aQ)

if, as a first approximation, the activities of A1 and Ti

are considered equal to their concentrations in weight

percent, in.this case 5.45w/o A1 and 8.4w/o Ti, the

dissolved equilibrium oxygen content at 1616°K (the

temperature at which this alloy solidified) is calculated

to be 1.7 x 10~ w/o for Equation 22 and 1.8 x 10~ w/o for

Equation 24. Therefore, for Ti02 to form in these alloys,

134

the concentration of dissolved titanium between the

dendrite arms must be approximately ten times higher

than the dissolved aluminum content. Thus, secondary

Ti02 inclusions did not form in these ternary alloys.

If all the dissolved equilibrium oxygen

(calculated from C + 0 = CO equilibrium) reacted with the

ternary alloys, the secondary inclusions would

- Sw occupy only 1.9 x 10 /o of the volume. The

photomicrographs of the ternary alloy microstructures,

as shown in section 6.4, verify that the volume occupied

by the secondary inclusions is very small. The primary

inclusions are much larger in size than the secondary

inclusions. The size of the primary inclusions vary in

size from 1 to 80 microns. They are usually found in the

dendrite arms, but they can also be present in the

interdendritic regions, since these primary inclusions can

be pushed by the thickening dendrites. Figure 29

shows a photomicrograph of the 93w/o Ni-5w/o Al-2w/o Ti

ternary alloy; note the large primary inclusions in the

dendrite. From this photomicrograph and the above

discussion, it can be seen that the contribution of

primary inclusions to the total inclusion volume is much

greater than that of secondary inclusions. Thus, the main

approach to reducing the total volume of inclusions present

in these alloys is to decrease the primary inclusion content.

135

This has to be achieved by restricting the dissociation of

the MgO and AÔg crucibles. The dissociation can be

avoided by increasing the vacuum pressure.

Spectrographs analysis of the castings indicated

little difference in the oxygen content between those

alloys melted in MgO and AÔ^ crucibles. For those

castings which were prepared by melting the alloy

constituents in AÔ^ crucibles, the average oxygen was

19.3 ppm. The castings which were prepared by melting the

alloy constituents in MgO crucibles had an average oxygen

content of 20.4 ppm. The oxygen content of the castings

obtained by spectrographic analysis was the total oxygen

present; this includes the oxygen present in the melt as

primary inclusions; as well as the dissolved oxygen.

Figure 56 shows the calculated C + 0 = C0(g)

relationships at CO partial pressures of 0.07 and 0.012 atm.

The oxygen and carbon concentrations, as obtained by

spectrographic analysis on the various castings investigated

in this study (see Table 4), are also shown in Figure 56 .

These plotted data points represent the total oxygen content,

the oxygen found in the primary inclusions, as well as the

dissolved oxygen. Thus, the actual dissolved oxygen

content in equilibrium with the carbon is lower than

these values.

136

e cx a.

c <u M >. K o

24

0.07 atm. CO

0.012 atm. CO

.002 .012 .004 . 006 .008 .010 / o C

Figure 56: The C + 0 = CO/ v relationship at CO partial pressures of 0. 07 and 0.012 atm -- Solid lines represent calculated values. Dots indicate values found in this investigation.

137

Table 8 contains the recommended vacuum pressures

for MgO and crucibles when they are used to melt

nickel-base alloys containing 0.003, 0.01 and 0.1w/o C.

For the nickel-base alloy containing 0.1w/o C, the pressure

above the melt must be greater than atmospheric, in order

to prevent dissociation of the MgO crucible; therefore,

alloys with C concentration greater than 0.01w/o should be

melted in AÔ^ at the recommended pressure given in

Table 8.

In summary, if nickel-base alloys of low carbon

contents are melted in MgO and AÔ^ crucibles under a O

vacuum less than 1 x 10 atm., the crucible will

dissociate. This dissociation leads to formation of

AÔg primary inclusions. In order to stop crucible

dissociation, the vacuum system pressure must be

increased. The volume content of secondary inclusions due

to microsegregation is negligible.

138

TABLE 8

RECOMMENDED MINIMUM VACUUM PRESSURES FOR Ni-BASE ALLOYS

Alloy Crucible

Ni-Base w/o C

MgO A12°3

Recommended pressure in torr

0.003

0 .010

0 .100

3

106

do not use (P>1 atm.)

9 x 10

7 x 10

7

-3

- 2

CHAPTER 7

CONCLUSIONS

As a result of this research program, the following

conclusions were reached:

(1) The composition of the Ni-Ti eutectic reaction,

£ t n+Y, is 12.6w/o Ti.

(2) The nickel-rich portion of the Ni-Ti-Al system

contains a ternary eutectic near the Ni-Ti binary. This

eutectic caused a trough to develop on the liquidus

surface. Across the 90w/o Ni vertical section, the trough

temperature minimum was at a composition of approximately

90w/o Ni-0.5w/o Al-9.5w/o Ti.

(3) The 90w/o Ni-Ti-Al alloy microstructure

generally consists of a y-y1 mixture between cored y

dendrite arms.

(4) The y-y1 mixture is a manifestation of

nonequilibrium solidification, since at equilibrium these

alloys should consist of the single phase, y.

(5) The microsegregation of titanium and

aluminum in these alloys is strongly effected by

composition. As a result of the trough in the liquidus,

alloy composition near the trough minimum composition

139

140

have segregation ratios of aluminum and titanium near

unity. By increasing the ^"/Ti ratio, the segregation

ratios of these solutes increases.

(6) Within the dendrite arms there is a greater

segregation of titanium than aluminum. For these alloys,

the maximum titanium SR experienced in the dendrite arms

was 1.68, with corresponding aluminum SR being 1.20.

(7) The solidification rate has little effect on

the segregation of titanium and aluminum.

(8) The dependence of secondary arm spacings on

local cooling rate was found to remain fairly constant for

the 90w/o Ni-Ti-Al alloys; this relationship is established

as d=27.7 (GR) 0-37 microns.

(9) The MgO and crucibles used to contain

these Ni-Ti-Al alloys dissociates during vacuum melting.

This dissociation of these refractories leads to oxygen

being pumped into the melt. This dissolved oxygen

subsequently reacts with the dissolved aluminum in the melt

and forms A^O^ inclusions. The dissolved oxygen does not

react with the dissolved titanium in the melt since the

deoxidation constant of aluminum in the nickel melt is

much greater than that of titanium.

(10) Those alloys melted in MgO, in addition to

AI2O3 inclusions, contain particles of MgO which reacts

with aluminum and oxygen in the melt to form Mg AloO/.

141

(11) The secondary inclusions in these alloys were

AI2O3. Although the titanium segregates to a greater

extent than aluminum in the interdendritic regions, the

aluminum still controls the allowable soluble oxygen

content.

(12) Primary inclusions comprises a much greater

percentage of the total inclusion content than does the

secondary inclusion.

(13) If MgO or is to be used to melt

nickel-base alloys, extremely low vacuum pressures should

not be employed since the crucibles will dissociate and

pump oxygen into the system with the end result being

formation of primary inclusions. Achieving better

cleanliness in these alloys will require attention to

these melting practice considerations.

APPENDIX A

THERMODYNAMIC DATA

The following paper has been submitted to the

Canadian Metallurgical Quarterly for publication. The

paper, including its references, is presented in its

entirety. The data in this paper was used to make all

the thermodynamic calculations which were required in this

solidification study.

142

143

The Thermodynamics o£ Dilute Liquid Nickel Alloys

G.K. Sigworth, J.F. Elliott, G. Vaughn, and G.H. Geiger

Abstract

The published data on the thermodynamics of liquid nickel

base alloys have been reviewed. Recommended thermodynamic values

are tabulated for binary and ternary alloys and calculated values

of deoxidation constants are given for selected elements.

Geoffrey K. Sigworth is an Assistant Professor in the Department of Metallurgy and Materials Science, Carnegie-Mellon University, Pittsburgh, Pennsylvania.

John F. Elliott is a Professor of Metallurgy at the Massachusetts Institute of Technology, Cambridge, Massachusetts.

Glen Vaughn is a graduate student in the Department of Metallurgical Engineering, University of Arizona, Tucson, Arizona.

Gordon H. Geiger is a Professor in the Department of Metallurgical Engineering, University of Arizona, Tucson, Arizona.

144

The Thermodynamics of Dilute Liquid Nickel Alloys

A good deal of information on the thermodynamic behavior of elements in

liquid nickel has been reported in the literature. Unfortunately, these data are often

widely scattered and presented in a variety of ways. Some systems have been

reviewed by Hultgren and co-workers^ in their survey of the thermodynamic

properties of binary metallic alloys. There is also thermodynamic information in

(2) the surveys on binary phase diagrams, but no single compilation has been made. In

this paper, the available thermodynamic data have been reviewed and summarized, and

recommended thermodynamic values are given for elements dissolved in nickel-based

alloys.

Three composition coordinates have been used in the literature reviewed

in this work: atom fraction (X), atom percent (a/o) and weight percent (%). For

simplicity, y is used herein to represent the activity coefficient when the pure substance

is used as the reference and standard states, and when atom fraction is used as the

composition coordinate. The symbol f is used to represent the activity coefficient

when the infinitely dilute solution is the reference state. It is to be noted that a

"hypothetical" standard state results when one uses the infinitely dilute solution as the

reference state. Unit activity at the "hypothetical" standard state Is obtained by the

relationship

at = f° • = 1; when = 1 (1)

Is the general composition coordinate and f? is the activity coefficient at infinite

dilution. It is to be recognized that the actual activity of i at composition = 1 is

not necessarily equal to 1, since the actual activity coefficient, 1, may no longer be

equal to one.

Although data in the literature appear in several forms, only two are used

In reporting the results of this study. They are composition in atom fraction with the

pure substance as the reference and standard states, and composition in weight percent

with the Infinitely dilute solution as the reference state and a hypothetical 1 percent

145

solution as the standard state. By convention, the activity coefficient Q./X^ =

when Xj - 0, and0./^ = f° = 1 when %. - 0. Since some publications do not employ

these standard and reference states, it has been necessary in many instances to make

a conversion from the data as reported in the literature. Should the reader wish to

alter the standard state or composition coordinate employed in this paper, a brief (3)

treatment of the method for making the conversion is available in an earlier compilation.

(4) A more generalized treatment of the subject is also available.

Table I shows the selected values for the standard free energy of solution

of elements in dilute liquid nickel. Tables II and III show selected Gibbs free energy

interaction coefficients for nickel-based alloys. The interaction coefficient was /ei

Introduced by Wagner, first used by Chipman, and later extended formally by

(7-12) (7) Lupis and Elliott. Using the notation of Lupis and Elliott,

Gf/RT = iny, =0ny? + E [X ] + E pjlX]2 (2) 1 1 1 j=2 3 j=2 J

+ Z 2 p{,k [X ] [Xfc] + 0 (X3)

J=2 k=2

j<k

The solvent, liquid nickel, has been designated as component 1 in the n-component

system, and the pure substance is used as the reference state. Third and higher

order terms are usually neglected, since the accuracy of the available data rarely

permit their calculation with any degree of certainty. When the composition coordinate

Is weight percent, the Taylor series expansion corresponding to eq. (2) is:

logf = S e| [wt.%.] + E rj [wt.% ] 2 1 j=2 1 ^ j=2 J

+ E E r|,k [wt.% j ] [wt.%k ]+ 0 (%3) (3)

j=2 k=2

j< k

Table I gives the selected values for the standard Gibbs free energy of

146

solution of elements in liquid nickel. Notes regarding the calculation of the

tabulated values appear with the table. Table II gives selected free energy inter

action coefficients for binary alloys, and Table III gives selected values for ternary

nickeHiased alloys.

In Tables n and in, Interaction coefficients obtained directly from

experimental measurements are italicized. (Others have been calculated by using

(7) the conversion equations of Lupis and Elliott. ) The temperature given indicates

the experimental temperature used in the original determination of thermodynamic

properties. A temperature range indicates that more than one temperature was

employed in the original experiments, and that the values tabulated are valid for

that range of temperature. The numbers of the references providing the results

shown in the tables are italicized. When the authors have found it necessary to

calculate (or recalculate) interaction coefficients from the dava given in a study,

an asterisk follows the corresponding reference number, •

Unfortunately, it is not possible to indicate In a straight-forward manner

the accuracy of calculations using the values tabulated in this study. Generally

speaking, errors tend to increase with increased concentration of the solute element.

It would be best, therefore, to consult the original works cited when second order

terms become appreciable at the compositions encountered in a calculation. All

references consulted are shown, and the ones used principally in determining the

tabulated data are italicized.

I Ik The cross product second order terms, pj' and rj* , have not been

tabulated in this study. These terms are generally obscured by the errors Inherent

in the measurement of the tabulated interaction coefficients. Even so, when the

inclusion of these terms is felt to be necessary, they often may be calculated from

(7) the reciprocal relations given by Lupls and Elliott

.{ = 2Pjl • I

Cl (5)

k l,k €. = p.' + J J

el = pl'j+ ej J Hk I

(6)

147

(7) and from conversion relationships/

A special note should be made regarding the values given for the interaction

coefficients in Ni-i-C alloys in Table III. The calculated interaction coefficients have

been determined for carbon-saturated solutions. They differ from the other coefficients

In that they are determined at a unit carbon activity; hence, X. 1. The free energy

(12) interaction coefficients are defined after Lupis:

*t 50ny € — — —_ ° 5Xi

and

* i 61og f

(V

T,p,ac = i

6%t (8)

T,p,ac = i

(12) Several conversion relationships are also given by Lupis, but the

most Important In the context of this paper is given below:

;• _ cc +Yc »cc C 13 (9)

1+X0c° + 2<XC)2P°

'The mole fraction of carbon at carbon saturation, X_, is known, but the

other Interaction coefficients are not known. It therefore has been

necessary to assume they are zero, or has been assumed to be equal to

ej . The calculated values in Table III reflect this assumption. These

values are considered to be rough estimates and are, therefore, shown in

parentheses.

The free energy data in Table I and information on the

thermochemistry of oxides have been used to calculate deoxidation

equilibria. The results are shown in Table IV.

148

REFERENCES

1. R. Hultgren, P. D. Desai, D. T. Hawkins, M. Gleiser and K. K. Kelly.

Selected Values of The Thermodynamic Properties of Binary Alloysf Amer.

Soc. Metals, Metals Park, Ohio. 1973.

2. (a) M. Hansen and K. Anderko. Constitution of Binary Alloys T

2nd. ed., McGraw-Hill, New York, 1958.

(b) R.P. Elliott. Constitution of Binary Alloys. First SupplementT

McGraw-Hill, New York, 1958.

(c) F. A. Shunk. Constitution of Binary Alloys. Third Supplement.

McGraw-Hill, New York, 1969.

3. G.K. Sigworth and J.F. Elliott. Can. Met. Quart. 13;455 (1974).

4. J. F. Elliott and J. Chipman. Proc. of Int. Svmp,, on Met. Chem.. Sheffield.

1971. p. 121. Iron Steel Institute, London, 1973.

5. C. Wagner. Thermodynamics of Alloys, p. 51, Addison-Wesley, Reading, Mass.,

1962.

6. J. Chipman. J. Iron Steel Institute. London. 180;97 (1955).

7. C.H.P. Lupis and J.F. Elliott. Acta Met. 14; 529; 1019 (1966).

8. C.H.P. Lupis and J.F. Elliott. J. Iron Steel Institute. London. 203:739 (1965).

9. C.H.P. Lupis and J.F. Elliott. Trans. Met. Soc. AIME,. 233:257;829 (1965).

10. C.H.P. Lupis and J.F. Elliott. Acta Met. 15:265 (1967).

11. C.H.P. Lupis. Liquid Metals. Chemistry and Physics. S. Beer, ed., pp. 1-36,

Marcel Dekker, New York, 1972.

12. C.H.P. Lupis. Met. Trans. 5:1919 (1974).

13. F. Vachet, P. Desre and E. Bonnier. Comptes Rendu Acad. Sci. Paris.

260:453; 1943 (19G5).

14. V.G. Dyubanov, A. Ya. Stomaxin and A.F. Filippov, Izv. Vyssh. Ucheb.

Zaved.. Chern. Met. 1972 f3):69.

15. V.V. Averin, P. A. Cherkasov, and A.M. Samarin. Issled. po Zharoprochn.

Splavam^ Akad. Nauk SSSR. Inst. Met. 9:204-18 (1962).

16. V.V. Averin, P.A. Cherkasov, and A.M. Samarin. Tr. Inst. Met, im.

A.A. Baikova. Akad. Nauk SSSR. No. 14:58-67 (1963).

149

17. V.V. Averin and' P.A.'Cherkasov. Fiz.-Khim. Osnovy Met. Protsessov. Komis. po

Fiz.~Khim. Osnovam Proizv. Stali. Sb. Statei.1964:79-85.

18. V.M. Sandakov, Yu. O. Esin, and P.V. Gel'd. Zh. Flz. Khim. 45:1798-99 (1971).

19. J. Golonka and J. Dziemianko. Prace Inst. Hutniczych. 21:267-273 (1969).

20. A. Ya. Stomakhin and A. Yu. Polyakov. Izv. Vyssh-Ucheb. Zaved.. Chern. Met.

1967 (3):116-121.

21. W. Oelsen and W. Middel. Mitt. Kaiser Wilhelm Inst. Eisenforsch. 19:1-26 (1937).

22. F. Korber, W. Oelsen, and H. Lichtenberg. ibid: 131-159.

23. A. Neckel, L. Erdelyi, G. Sodeck, and E. Buschmann. Can. Met. Quart. 13:

379-381 (1974).

24. A.K. Biswas and E.J. Grimsey. Inst. Min. Met.. Trans. 83:250C-255C (1974).

25. Yu. O.Esin., V.M. Baev, P.V. Gel'd, and M.S. Petrushvsky. Izv. Akad.Nauk

SSSRf Metally. 1974 (4):73-76.

26. Y. Kojima and K. Sano. Tetsu-to-Hagane. 47:897-902 (196H.

27. H. Schenck, M.G. Frohberg, and E. Steinmetz. Arch. Eisenhuetenw. 34:37-42

(1963).

28. L.D. Lucas. Mem. Sci. Rev. Met. 66:747-755 (1969).

29. E.T. Turkdogan, R.A. Hancock, and S.I. Herlitz. J. Iron Steel Inst.. London.

182:274-277 (1956).

30. T. Fuwa and J. Chipman. Tetsu-to-Hagane. 45:600-607 (1959).

31. V.I. Artemov, V.T. Burtsev, V.I. Kashin, and L.N. Saksonova. Izv. Akad. Nauk

SSSR. Metally. 1972 (4):25-31.

32. K.A. Matsarin, S.V. Kraskovskly, and F. P. Edneral. Izv. Vyssh. Ucheb. Zaved..

Chern. Met. 1966 (3):68-72.

33. G.K. Sigworth and J. F. Elliott. Metal Science. 8:298-310 (1974).

34. N. MeyssonandA. Rist. Rev. Met. 62:1127-1131 (1965).

35. L. Sh. Tsemekhan, S. E. Vaisburd, and Z. F. Shirokova. Zh. Fiz. Khim. 45:

2074-2076 (1971).

36. J. Vrestal and J. Kucera. Trans. Met. Soc. AIME. 245:1891-1895 (1969).

37. J. Vrestal and J. Kucera. Met. Trans. 2:3367-3372 a971).

150

38. B. Predel and R. Mohs. Arch. Elsenhuetenw. 41:61-66 (1970).

39. K. W. Lange and H. Schenck. Met. Trans. 1:2036-2038 (1970).

40. B.F. Belov and I.A. Novokhatsky, Izv. Akad. Nauk SSSR. Metallv. 1970

(51:69-72.

41. C.B. Bezobrazov and A.M. Samarin. Izv. Akad. Nauk. Otdel. Tectu Nauk.

1953 (12):1790-1796.

42. R.J. Fruehan. Trans. Met. Soc. AIME. 242:2007-2008 (1968).

43. S.W. Gllby andG.R. St. Pierre. Trans. Met. Soc. AIME. 245: 1749-1758 (1969).

44. S. Wagner, D.J. Shade, andG.R. St. Pierre. Met. Trans. 3: 47-50 (1972).

45. R. Hultgren, R. Orr, P. Anderson, and K. Kelly. Selected Values of Thermodynamic

Properties of Metals and Alloys. J. Wiley and Sons, New York, 1963 (Revised

Supplement of October, 1966).

46. B. Predel and R. Mohs. Arch. Eisenhuetenw. 42:575-579 (1971).

47. Y. Tozaki, Y. Iguchi, S. Ban-ya, and T. Fuwa. Proc. Internat. Symp. Chemical

Metallurgy of Iron and Steel. July. 1971. pp. 130-132, 1973.

48. H.O. von Samson-Himmelstjerna. Z. Metallk. 28:197-202 (1936).

49. C.W. Schultz, G.R. Zellars, S.L. Payne, andE.F. Foerster, U.S. Bureau Mines

Rept. Invest. 6410, 1964. (As given in reference 1.)

50. A. El-Kasan, K. Abdel-Aziz, A.A. Vertman, and A.M. Samarin. Izv. Akad. Nauk

SSSR. Metally. 1966 (3):19-30.

51. A.D. Kulkarni and R.E. Johnson. Met. Trans. 4:1723-1727 (1973).

52. G.R. Belton and R. J. Fruehan. J. Phys. Chem. 71:1403-1409 (1967).

53. R.J. Fruehan, Ph.D. Thesis, U. Pennsylvania, 1966.

54. K.C. Mills, K. Kinoshita, and P. Grieveson. J. Chem. Thermodyii. 4:581-590 (1972).

55. G.R. Zellars, S.L. Payne, J.P. Morris, andR.L. Kipp. Trans. Met. Soc. AIME.

215:181-185 (1959).

56. R. Speiser, A.J. Jacobs, andJ.W. Spretnak. ibid.. 185-192.

57. A.P. Lyubimov, V. Ya. Zobens, and V.I. Rokhovskiy. Zh. Fiz. Khim. 32:1804-1808

(1958).

58. B. Predel and R. Mohs. Arch. Eisenhuetenw. 41:143-149 (1970).

59. G.I. Batolin, N.N. Minenko, andV.S. Sudavtsova. Izv. Akad. Nauk SSSR. Metally.

1974 (5):99-103.

151

60. V.D. Ivanova and B. P. Buiylev. Zh. FLz.Khim. 45:2056-2057 (1971).

61. A. Sleverts and W. Krumbhaar. Ber. Deut. Chem. Ges. 43:893-900 (1910).

62. T. Bags haw and A. Mitchell. J. Iron Steel Inst.. London. 204:87-90 (1966).

63. T. Bagshaw, D. Engledow, and A. Mitchell. J. Iron Steel Inst. . London.,

203:160-165 (1965).

64. T. Busch and R. A. Dodd. Trans. Met. Soc. AIME. 218:488-490 (1960).

65. E.S. Levin, T.K. Kostina, M.S. Petrushevsky, P.V. Gel'd, andK.T. Kurochkin.

Izv. Akad. Nauk SSSR. Metally. 1973 (31:67-72.

66. T.K. Kostina, B.A. Baum, P.V. Gel'd, andK.T. Kurochkin. Izv. Vyssch.

Ucheb. Zaved.^Tsvet. Met. 1970:22-27.

67. P.V. Gel'd, T.K. Kostina, B.A. Baum, andK.T. Kurochkin. Zh. Flz. Khim. 45:

1999-2001 (1971).

68. H. Schenck and K.W. Lange. Arch. Eisenhuetenw. 39:673-676 (1968).

69. M. Weinstein and J.F. Elliott. Trans. Met. Soc. AIME. 227:285-286 (1963).

70. R.G. Blossey and R.D. Pehlke. Met. Trans. 2:3157-3161 (1971).

71. N.N. Kalinyuk and V.I. Lakomskiy. Automat Weld. 20:27-32 (1967).

72. K.W. Lange, D. Ksinsik, and H. Schenck. Erzmetall. 25:232-238 (1972).

73. V.I. Lakomskiy and N.N. Kalinyuk. Izv. Akad. Nault SSSR. Metally. 1966 (2): 149-155.

74. H. Schenck and K.W. Lange. Arch. Eisenhuetenw. 37:739-748 (1966).

75. H. Wada, K. Gunji, andT. Wada. Trans. Iron Steel Inst.. Jap. 8:329-336 (1968).

76. E.G. Ivanov, A. Ya. Stomaxin, G.M. Medvedeva, andA.F. Filippov. Izv.

Vyssch-Ucheb. Zaved.. Chern. Met. 1966 (5):69-72.

77. J. C. Humbert and J.F. Elliott. Trans. Met. Soc,, AIME. 218:1076-1088 (1960).

78. R.G. Blossey and R.D. Pehlke. Trans. Met. Soc. AIME. 236:566-569 (1966).

79. A. Ya. Stomaxin, P. Bayer, and A. Yu. Polyakov. Izv. Akad. Nauk SSSR. Metally.

1965 (4):37-45.

80. H. Schenck, M.G. Frohberg, and H. Graf. Arch. Eisenheutenw. 30:533-537 (1959),

81. Y. Kojima, Y. Yamada, and M. Inouye. Tetsu-to-Hagane. 61:195-201 (1975).

82. Yu. M. Pomarin, G.M. Grigorenko, V.I. Lakomsky, G.F. Torkhov, and

A.V. Sherevera. Izv. Akad. Nauk SSSR, Metallv. 1972(4):32-36.

83. E.S. Tankins. Met. Trans. 1:1897-1904 (19701.

152

84. E.S. Tankins, J.F. Erthal, andM.K. Thomas, Jr. J. Electrochem. Soc.

112:446-450 (1965).

85. E.S. Tankins. Met. Trans. 1: 1465-1467 (1970).

86. E.S. Tankins and N.A. Gokcen. Met. Trans. 2:1605-1611 (1971).

87. E.S. Tankins. Met. Trans. 1:2637-2638 (1970).

88. E.S. Tankins, N.A. Gokcen, and G.JR. Belton. Trans. Met. Soc. AIME. 230:

820-827(1964).

89. H.A. Wriedt and J. Chipman. Trans. Met. Soc. AIME, 206:1195-1200 (1956).

90. W.A. Fischer, D. Janke, and W. Ackermann. Arch. Eisenhuetenw.41:361-367 (1970).

91. W.A. Fischer and W. Ackermann. Arch. Eisenhuetenw. 37:43-47 (1966).

92. V.V. Averin, A. Yu. Polyakov, and A.M. Samarin. Izv. Akad Nauk SSSR. Otdel.

Tekn. Nauk. 1957(8): 120-122.

93. W.A. Fischer and W. Ackermann. Arch. Eisenhuetenw. 37:779-781 (1966).

94. J.E. Bowers. J. Inst. Metals. 90:321»328(1962).

95. B.F. Belov and I. A. Novokhatskiy and Yu. A. Lobanov. Izv. Akad. Nauk

SSSR. Me tally. 1967 (3):53-62.

96. W.A. Fischer and D. Janke. Z. Metallic. 62:747-751.

97. N.A. Vatolln, A.I. Tlmofeyev, O.A. Yesln, andE.L. Dubinin. Zh. Fiz. Khtm. 45:

2021-2026 (1971).

98. N.A. Vatolin, A.I. Timofeyev, andE.L. Dubinin, ibid.. 2027-2029.

99. W.V. Venal andG.H. Geiger. Met. Trans. 4:2567-2573 (1973).

100. S.E. Vaisburd, V.G. Dyubanov, I.N. Zedina.A. Ya. Stomaxin, andA.F. Fillppov.

Izv. Vyssch. Ucheb. Zaved.. Chern. Met. 1972(l):53-55.

101. J.A. Cordier and J. Chipman. Trans. Met. Soc. AIME. 203:905-907 (1955).

102. C.B. Alcock and F.D. Richardson. Acta Met. 6:385-395 (1958).

103. C.B. Alcock and L.L. Cheng. J. Iron Steel Inst.. London. 195:169-173 (1960).

104. V. Ya. Dashevskiy and A. Yu. Polyakov. Izv. Akad. Nauk SSSR. Metally.

1966(5):34-41.

105. V. Ya. Dashevskiy and A. Yu. Polyakov. Fiz.-Khim. Osn. Proizvod. Stali. 1968:

49-52.

106. G.A. Meyer, J.S. Warner, Y.K. Rao, and H.H. Kellogg. Met. Trans. 6B-.229-235

(1975).

153

107. P.A. Cherkasov, V.V. Averin, and A. M. Samarin. Izv. Akad. Nauk SSSR.

Metally. 1967(l):49-55.

108. P.A. Cherkasov, V.V. Averin, and A.M. Samarin. Zh. Fiz. Khim. 42:767-771

(1968).

109. P.J. Bowles, H.F. Ramstad, and F.D. Richardson. J. Iron Steel Inst..

London. 202:113-121 (1964).

110. Yu. V. Tarakanov, P.A. Cherkasov, V.V. Averin, and A.M. Samarin.

Dokl. Akad. Nauk SSSR. 163:166-168 (1965).

111. K. Schwerdtfeger and H.-J. Engell. Trans. Met. Soc. AIME. 233:1327-1332 (1965).

112. Yu. M. Gertmann and P.V. Gel'd. Tr. Uralsk. Politekhn. Inst. 1961 (114):96-106.

113. P.V. Gel'd, and Yu. M. Gertman. Fiz. Metal. Metalloved.. Akad. Nauk SSSR.

10:299-300 (1960).

114. L„ Martin-Garin, I. Ansara, and P. Desre. Compt. Rend.. Sec. C. 266: 1595-1597

(1968).

115. V.N. Eremenko, G.M. Lukashenko, andV.L. Pritula. Zh. Fiz. Khim. 45:

1993-1995 (1971).

116. N.R. FrageandYu. G. Gurevich. Izv. Vyssch. Ucheb. Zaved. SSSR. Chern. Met.

1974 (8):5-8.

117. A. Ya. Stomaxin and A. Yu. Polyakov. Izv. Akad. Nauk SSSR. Metally. 1967 (2):

49-54.

118. W. C. Ballamy and E.E. Hucke. J. Metals. 22(8):43-50 (1970).

119. V.I. Fedorchenko, V.V. Averin, and A.M. Samarin. Dokl. Akad. Nauk SSSR.

196:1093-1096 (1971).

120. V.I. Fedorchenko, V.V. Averin, and A.M. Samarin. Izv. Akad. Nauk SSSR.

Metally. 1971 (3V.73-77.

121. W. Flurschuetz. Abhandl. Deut. Akad. Wiss. Berlin. Kl. Math.. Physik. Tech.

1962. No. l! 385-389.

122. R.M. German and G.R. St. Pierre. Met. Trans. 3: 2819-2823 (1972).

123. K.W. Lange. Z. Metallkunde. 64:111-112 (1973).

124. A.I. Tripolitov, O.N. Kryuchkov, A. Ya. Stomaxin, andA.F. Fllippov. Izv.

Vyssch.Ucheb. Zaved.. Chern. Met. 1969 (5):57-61.

154

125. A.L. Marchant, R.M. Prado, J. Fava and R.A. Flinn. Trans. Amer. Found.

Soc. 80: 193-196 (1972).

126. V.P. Fedorchenko, V.V. Averin, and A.M. Samarin. Dokl. Akad. NaukSSSR.

183: 894-6 (1968).

127. K.W. Lange and H. Lindscheid, Proc. Internat. Symp. Chemical Metallurgy of

Iron and Steel, Sheffield, July, 1971, pp. 133-136, 1973.

128. P. Sieben and N.G. Schmahl. Giesserei, Tech.-Wiss. Beih. Giessereiw. Metallk.

J8: 197-211 (1966).

129. T.N. Turner and G.H. Geiger. Proc. Electric Furnace Conf., 1968. pp. 114-118,

AIME, New York.

130. T. Fuwa, M. Fujikura, and S. Matoba. Tetsu-to-Hagane. 46:235-237 (1960).

131. K.W. Lange and H. Schenck. Z. Metallk. 60:638-642 (1969).

132. T.V. Lipets, Zh. L. Vert and I.P. Tverdovskiy. Zh. Fiz. Khim. 43: 1841-1844

(1969).

133. K.W. Lange and H. Schenck. Arch. Eisenhuetenw. 40: 737-741 (1969).

134. K.W. Lange and H. Schenck. Z. Metallk. 60:62-68 (1969).

135. T. Saito. Sci. Rep. Res. Inst., Tohoku Univ., Series A.1:419-424 (1949).

136. H. Sakao and K. Sano. J. Jap. Inst. Met. 26:30-38 (1962).

137. H. Sakao and K. Sano. ibid.: 236-240.

138. W.A. Fischer and D. Janke. Z. Metallk. 62:747-751 (1971).

139. D. Janke and W.A. Fischer. Arch. Eisenhuetenw. 44:15-18 (1973).

140. P.A. Cherkasov, V.V. Averin, and A.M. Samarin. Izv. Akad. Nauk SSSR,

Metal!y. 1967 (l):49-55.

141. V. Ya. Dashevskiy and A. Yu. Polyakov. Fiz.-Khim. Osnovy Proizv. Stali.

1968:49-52.

142. J.J. deBarbadillo, paper presented at Amer. Vacuum Soc. Conf., June 23, 1975,

Columbus, Ohio. (Available from Amer. Vacuum Soc. on microfiche.)

143. J.J. deBarbadillo, private communication.

144. K. Goto, S. Ban-ya, and S. Matoba. Tetsu-to-Hagane Overseas. 3_:184-189

(1963).

145. J.F. Elliott and M. Gleiser. Thermochemistry for Steelmakinq, Vol. I.

Addison-Wesley, Reading, Mass., 1960.

Table I. The Gibbs Free Energy of Solution of Elements In Liquid Nickel

Element, 1

(a)

<1600° C)

AG° (X)

(b)

(cal/gm. -atom)

AG°$)

(b)

(cal/gm. -atom)

T(°C)

(c)

References Consulted

Al(l)

Au{l)

B(s)

Ba>

C(s)

Ca(l)

Ca (g)

Co(l)

0.00025

1.62 0.0016 (0.009)

0.31 ( 0 . 6 ) 1.27

0.45

-37000 +3.31 T

1800 (-14200 - 5.22T)

(-26200) '

4980-,5i01 T - 1920

-44520 +24.25T

300 - 1.74T

(1)

(1)

-37000 " 4.28T 1800 -11.54T

(-14200-11.0 T\( '

(-26200 -5. 78)( '

(1) ' 4980-10.S9 T - 1920 - 8.37T

-44520 + 15.88T

300 - 10.88T

( 1 )

1600 13-18,19,20-22

1460 1,23,24

1600 25

1600 25

1560 15,26-32,125,144 1477 1, 34

1477 1, 34

1600 35-37,38,39, 127

Cr (s)

Cr(l)

Cu(l)

Fe (1)

0.46

(0.39)

2.18 0.36

2500 - 2.86T

-2200 - 0.69T

2900

-10000 + 3.28T

2500 - 11.85T

-2200 - 9.58T

2900 - 9.29T

-10000 - 5.75T

1550-1600 15,17,40-42,43^-45

1550-1600 15.17,40-42,43-45

1550 50, 51

1510-1700 _1,22,35,3 8,43, 47,

50,52-56,57-59

Gea) (9)

Mgfl)

Mg(g)

0.13

(0.32)

2.2

- 7470

( - 4200) (1)

(-38910 + 22.34T) (D

-7470 - 9.55T

4800 + 8.36T

(-4200 - 7.38T)'

(-38910+ 14.96T)

(!)

(1)

1450 JS0

1450-1700 _61-^9, 70-74

650-850

650-850

128 128

Mna) Mo(s)

Moa) |N2(g)

1 2.1 (1)

(0) (1)

(7780 - 2.69T)

(0)1

(1) ( - 9.0T)

(7780 - 12.79T)(1)

( - 10. lT 1* 6660 + 9.49T

1600 16,_17

2_ 2

1550-2200 39,75-J79, 80-82, 126 Ui Ui

Table I. The Gibbs Free Energy of Solution of Elements in Liquid Nickel

Element, i

(a)

(1600°C)

o (b) AGt (X)

(cal/gm. -atom)

AG° (%)

<b)

(cal/gm. -atom)

T°(C)

(o)

References Consulted

|o2(g)

Pb (1)

Pd(s) f S ( g )

Si(l)

1.4

1.35

0.00014

(1200)(1)

1120

-42000 + 4.83T

t16,970 + 0.336T

(1200 - 11.63T)

1120 - 10.31T

- 28340 +3.62T - 42000 - 2.84T

1500 - 1700

1600

1500 - 1600

1500 - 1650 1550 - 1600

83-90,91,92,93,94, 95.96,136

50

1. 91. 99-101. 102-106 15,17721,25.50,

107-110, 111-114

Sn(l)

Ti(s)

V (s)

V(l)

0.14

0.00019

0.011 (0.009)

(-25000 + 9.46T1

(-28300 - 1.93T)(2)

(-12450.- 2.3T)( '

(-17600)'

(-25000 - 1.07T)

(-28300 - 10.66T)

(-12450 - 11.15T)'

(-17600 - 8.85T)

< 2 ) (1)

(1)

1300 _1, 25, J15

1550 -1700 15-17.20.107.116.117-322

1600 15

1600 15

W(s)

W(l)

Zr (s)

13.5

(U.3)

0.00007

(9000)^' '

(-69000 + 17.87T)

(16500 - s.esT)* ,A3Jf

(16500 - 15.04T)(1'3-

( 9000 - 11.39T>(1,3^

(-69000 + 7.87T) <2>

1600 121, 123

1600 121, 123 1550 -1700 20. 124

Notes: (a) For the Standard State shown in the first column. Gases are at one atmosphere pressure. Values in

parentheses are for unstable standard states at 1600°C.

(b) See the text for an explanation of the standard and reference states used. Values in parentheses

are considered to be uncertain.

(c) Principal references used in data selection are italicized.

(1) Regular solution assumed for the liquid phase.

(2) Calculated from nitride solubility.

(3) Calculated from carbide solubility. I-1 t_n CT>

Table n. Free Energy Interaction Coefficients In Binary Nickel Alloys

. Element, 1 o

T ( C) References

A1

Au

Ca

Co

(2*) <-M)

(L.) <£>

(2J

(M)

(OJ

(0)

(0.08) (0. 003)

(0. 004)

(0)

(-0.0006) (0)

(0) (0)

1600

1460

1477

1600

13

23*

J*. §£* * 35,36,127

Cr 1^8 (-1J 0.0083 (0) 1600 42,43,44 Cu 1x1 (0) 0.0076 (0) 1600 i,49~" Fe -^1-2 JL 0.013 o 1510 - 1600 1*43.55*.56*.57 H 1«° 0.5 0 jQ_ 1500 - 2400

N °«8 0.3 0 0 1500- 1700

61.65.70.71.73

79f82.120.126 1500 - 1700 84.88.89.92.94 ° 0.7 0.3 0 0

S -182000/T + 94.2 -82870/T + 42.8 -1453/T + 0.748 0 1500 -1600 go' "i of^inV i<u

« ML- (0) 0.11 . (-0.0-013) 1580 - 1610 m, m

Ln

Table III - Free Energy Interaction Coefficients in Ternary Nickel Alloys

Element, 1 * i eC

*1 ec

Part A. Nl-t-C Alloys

T (°C) References

A1 3.44 0.027 (3.44) (0. 027) (0.056) 1600 27 As 4.63 0.017 (4.63) (0.017) (0.081) 1600 27 Au 0.96 0.004 (0.96) (0. 004) (0. 003) 1600 27 B 3.53 0.064 (3.53) (0. 064) (0. 058) 1600 27 Ce - 4.92 - 0.006 (-4.92) (-0.006) (-0.12) 1600 27 Co - 0.29 - 0.001 (-0.29) (-0.001) (-0.023) 1400-1600 26.27 Cr - 2.54 - 0.013 (-2.54) (-0.013) (-0.071) 1400-1600 26.27

Cu

Fe

Ga

Ge

In

Mn

Mo

0.96.

0.96

2.97

5.80

2.97

0.44

- 2 . 6 2

ta) 0,004

0,004

0.012 0.021 0.009

0.0017

- 0.005

(a) (0.96 V 0.8^a)

(2.97)

(5.8)

(2.97)

(0.44)

(-2.62)

,0,004,

(0.012) (0.021) (0. 009)

(0.0017)

(-0. 005)

(0. 003) 0<a)

(0. 046)

(0.11) (0.05)

(-0.008) (-0.073)

1600 1560-1600

1600

1600

1600

1600

1400-1600

27. 125

27, 3(£, 130*

27

27

27

26,27

26. 27

P

Pd

Pt

Sb

Se

Si

5.23

0

_0_

4.88

2.92

3.93

0.04

0.002 0.003 0.012 0.011 0.031

(5.23)

( 0)

( 0)

(4.88)

(2.92)

(3.93)

(0.04)

(0.002) (0.003)

(0.012) (0.011) (0^031)

(0. 094)

(-0.017)

(-0.017)

(0.087)

(0.045)

(0.066)

1600 1600 1600 1600 1600 1600

27

27

27

27

27

26,27 Ui 00

Table III (continued)

Sn 3.84 0. 01 (3.84) (0.01) (0.065) 1600 27 Te 2.92 0.008 (2.92) (0.008) (0.045) 1600 27

Tl -4,0 - 0. 022 H.0) (-0.022) (-0.10) 1600 27 V - 2.96 - 0.015 (-2.96) (-0.015) (-0.08) 1600 26.27

W - 2.62 - 0.001 (-2.62) (-0.001) (-0.07) 1600 27

Zn 1,92 0. 008 ( 1.92) (0. 008) (0. 024) 1600 27

Notes: (a) The values for dilute solutions have been taken from dat^ refergjnee 144»and those for carbon saturation are taken from data in references 27 and 130, The value of r / * .

V C

Part B. Ni-i-Ca Alloys

Element i Ca Ca Ca Ca

Ca e; HjDi references

Cr

Fe

Mn Mo

12.

7.7 8.7

32.

0

0

0 0

0.059 0.035 0.040 0.086

0

0

0

0.075 0.047 0. 053 0.20

1600 1500-1600

1600 1600

* 111 34 .142?

Part C. Nt-l-Cr Alloys

Element^ l Cr Cr Cr

l CCr

Cr 3l T f°C) references

Fe

Si

Tl

2 . 2 (10.) 11.3

0.01 (0.09) 0. 06

0.-91 (0. 05) 0. 055

1600 1600 1600

43, 44 +

107, 108 107. Hi*. 140 UI

VO

PartD. Nl-l-H Alloys

Element, i H H H

: 1 H

A1

Au

Co

Cr

Cu

Fe

2.0 3.54

0.72

0.84

0.33

0 .6

0.014

0.0076

0.0031 0.0036

0.0017

0,0024

- 0. 0002 0

0 - 0.0001

(0 ) 0

0.26 0.65

- 0.07

- 0.04

- 0.16 - 0.1

Mo

Mn

Si

V

W

3.36

- 2 .0 4.14

2.73

5.75

0

0

0

0.4

0

0.011 0.0096

0.033

0.013

0.011

0

0

- 0.0004

0.

0.0001

0.60 - 0.75

- 0.80 0.44

1.20

Part E. Nl-Fe-Mg Alloys

Fe

* Mg

6.6

Fe

(L^SL

(1.)

Fe 6 Mg

0.03

Fe f Mg

(0)

eM*

Fe

0.063

Element, t

A1

Ce Co Cr

Fe

N

0.54

-304.

-1.25

-22. - 4.3

•L 0.1

420. 0

-2.5

- 0 . 2

e N

Part F. Nl-l-N Alloys

_0_

• 0.55 • 0.0054

0.11 0.02

N

0

0 _0_

0

0

_N I

e

- 0.004

- 5.5 - 0.04

- 0.42

- 0.09

references

1500 - 1700

1600 1500 - 1700

1600

1500 - 1600

1600

65, 70, 131_

131

39, 62, 64, 68, 70*. 74. 133 62, 133*

* % •

62., 63, 70, 74

62,63,^i*,70*, 74, 133

1500 - 1700

1600

1460 - 1560

1600

1500 - 1700

134

J70 66. .67, 68

62

134

T(°C)

1600

references

142*. 143*

T (°C) references

1550 76, 117_ . 1550 J76

1600 39*, 78

1550 - 1600 H, 11 g

1600 - 2200 75*, J77, _78," 80-82

Mo -16. 9.8 -0.04 0 -0.3 1550 - 1600 79, 120, 126

Tl -37. -6.9 (- 0.20) 0 - 0.67 1550 - 1600 117. 120

W -189. 402. - 0.26 0 -3.46 1550 79, 120, 126

Zr - 86. 48. - 0.24 0 - 1.59 1600 124

PartG. Nt-t-O Alloys

I 1 1 I o

Element, i lo <>0 fo_

r_o li_ T(°C)

Au -2.4 28.7 0 0. 0001 - 0.05 1550

C (- 26.) - (-0.57) - - 0.43 1950

Co - 1.4 0.8 -0.006 0 - 0.034 1600

Cr -40.7 - 4.6 -0.20 0 - 0.66 1600

Cu - 2.1 1.5 -0,008 0.0001 - 0.045 1500 - 1600

Fe -6.4 2.9 - 0. 029 0.0002 - 0.11 1600

Mn -97. - 5.0 - 0.45 0 - 1.53 1600

P 0.5 0 _0 0 - 0.019 1600

S - 10.8 - 5.0 - 0.089 0 - 0.16 1600

SI - 14.6 -7.7 - 0.137 0 - 0.22 1600

Tl -86. - (-0.46) - - 1.37 1600

V -80. -9. -0.4 0 - 1.26 1600

references

* 87

• 31

85 , 90, 136 „

15, 17, 41*, 83, 137

84*, 96*

is. VL &*.&*, aa.*,9o, — J3\ I2£, 144

15, 1T_

139

_93

15. 17. 110

107. 140

15, 17

Part H.

it 11

Element. I fs ?S fs V

A1 14.6 7.8 0.133 0

Co lj_6_ 0 0.007 0

Cr 6.2 0.7 0.03 0

Cu 0 0 0.0003 0

Fe 1.1 0 0.005

Mo 19.5 -43. 0.053 -

SI 5.8 2.9 0.048

Tl 30. 5.6 0.16

0

0

Nl-t-S Alloys

references

0.11 1600 99, 105

0.009 1540 103*

0.046 1600 _99

- 0.004 1600 1 02

0.005 1540 - 1600 99. 101. 103

0.15 1600 99*

0.043 1540 104

0.24 1600 99

K>

163

Table IV. Calculated Deoxidation Equilibria in Nickel Alloys.

Deoxidation reaction log K K1873°K

2 A1+ 30 = A1203(S) 60,760/T - 18.7 5.49 x 1013

2 B + 30 = B203(JL) 47,383/T - 15.64 4.55 x 109

C + 0 = C 0 ( g ) 3 , 2 3 0 / T + 2 . 2 6 9 . 6 5 x 1 0 3

Ca + 0 = CaO(s) 27,706/T - 6.57 1.67 x 108

Co + 0= CoO(s) 9,815/T - 6.9 T 0.022

2 Cr + 30 = Cr203(s) 49,370/T - 18.51 7.03 x 107

Fe + 0 FeO(Jl) 6,593/T - 3.73 0.61

Hg. + 0 = MgO(s) 26,324/T - 7.57 3.07 x 106

Mn + 0 = MnO(s) 17,609/T - 6.52 7.61 x 102

Si + 20 = Si02(s) 32,980/T - 10.88 5.35 x 106

II + 20= Ti02(s) 35,540/T - 11.46 3.27 x 107

2J/ +30= V 20 3(s) 47,210/T - 17.04 1.46 x 10 8

Z r + 2 0 = Z r 0 2 ( s ) 3 4 , 0 0 0 / T - 7 . 5 2 4 . 3 0 x 1 0 10

APPENDIX B

DESCRIPTION OF ZAF PROGRAM

As is customary, the expression used to relate

observed intensity ratios (k^) to concentrations (C^) is

written as:

C. = (Z. A. F.) k. i. K i i x' 1

where (atomic number correction), A^ (absorption), and

F^(secondary fluorescence) are computed correction factors

for element i.

Since each of these three correction terms involve

knowledge of the concentrations for all elements in the

matrix, it is necessary to solve for the ZAF corrections in

an iterative fashion.

The iterative procedure used is the same as that

used by J. Colby (47) in the well-known Magic IV program:

1) Initially, the estimated concentrations are set

equal to the input k ratios;

c: = ki

2) The concentrations are normalized to sum to 100%.

C| c = _±

Z C! i x

3) Using the normalized C^ values, the ZAF factors

for each element are computed.

164

165

4) Using the computed ZAF factors and the

unnormalized previous concentration estimates

(C!), we compute

Ci lr ! = i 1 ZiAiFi

5) In order to generate a new estimate of concentra

tion, use is made of the observation that the

curve relating concentration and k-ratio for an

element is empirically found to have the form

l-k± 1-Cjl -j—- = CX. -p-t ki 1 ci

and thus, for our new estimate we also expect

i-k! i-c:

eliminating the constant CX^, we then estimate a

new set of concentrations:

Ci " k. (C[-k!)+k!(l-Cl)

6) These new estimates are compared to the previous

estimates (C^) and if |- C^|<.001 for all

elements, the iterations are terminated. Other

wise, all are set equal to the above computed

Ci and steps 2-6 are repeated.

It is necessary to compute the ZAF factors. To

compute the atomic number correction, the Duncumb-Reed (48)

expression for Z is used:

where F? = R../S.. i 11 11

166

F? Z. = — x P:

F C.R.. F: = UJJ. 1 T" C.S. .

J J i]

where R. . (backscatter coefficient) is a function of W.

w Eci

and Z_^ (the atomic number) . As discussed by Beaman and

Isasi (49) a polymonial expansion for R is used. The mean

stopping power (S) is obtained from:

ffy = 111(1.166^/^)

where

is the average excitation energy for the analyzed element,

A- is the atomic weight, E . is the critical absorption j CI

energy, and is the mean excitation potential for the

retarding element in the matrix. The values are computed

from a least-squares fit expression to experimental data of

Duncumb and Reed (48).

For the absorption correction the expression commonly

used is: fi(X)°

Ai = f±(X)'

ko _ 1+hi fi(x) =[1+|i][i+hi(i+fi)]

167

nV

f i (X) ' = £l+xj 1+h' ( 1+X± )J

1.2A. where h^ =

zi

and h' = T C.h. j J J

4.5 x 105

°i " ~r65~"~l765 o ci

X± = ( y / p)i;LCsc V

X! = C. ( y / p)..Csc * 3 3 3

where (y/p)ij

is the mass absorption coefficient of element j for the

analyzed radiation of element i. The (p/p)^ values are

computed from

nj •j K

where ^

is the wavelength of the analyzed radiation. The 13. and n^

constants are computed for least-squares fit expressions to

B and n values for energies greater than the N absorption

edge.

The program checks for secondary fluorescence by

comparing the absorption edge of the analyzed element

168

against the energy of the analyzed emissions for all

other elements. If fluorescence can occur, the F factor

is computed by the method of Duncumb and Reed (48)

F. = 1 f 1+2 G- -C. (X. . + Y . . ) ! L . l j J i j i r J

where the sum is over all analyzed lines capable of

fluorescing the ith element.

1.

ti l A j

i/j KCX LCX

KCX 1 4.2

L(X 2.4 1

E o i cj

E o 11 LEci

1.67

where -

r^ is the "jump ratio" for the Ec absorption edge and is

computed using the expression for

ri - 1 r. I

used by J. Colby in Magic IV (47).

w. is the fluorescent yield and is computed from an

expression of the form

[i ] * • aO+aIz+a3z3

169

For KCX radiation, the coefficients AQ> and A3 are

taken from a recent least squares fit by Bambynek (50),

for LCX radiation the coefficients are those used in

Colby's Magic IV (47).

1+ Xi Xij ln X./Csc V

XL

1+ gj

Yij ~ ln Xj/Csc v

Where X' and a are as previously defined above.

APPENDIX C

MICROSEGRE GATION DATA

The following data in Tables C-l through C-20

represent only a small portion of the microsegregation

data obtained during this investigation. These data are

provided for those who wish to make concentration maps

across secondary dendrite arms. Each Table contains a

schematic dendrite which indicates the location and

direction of the traverse. The selected numbers on the

dendrite arms correspond to the numbers in the Table and

give the relative position of the analysis.

170

171

TABLE C-l

MICROSEGKEGATION DATA FOR SAMPLE #1: DENDRITE A

Secondary Arm Spacing: 20\I

Starting Position of Point #1:

X = 23330, Y = 38464y

Point # Position,y Composition, a/o

X Y Ni A1 Ti

1 23330 38464 92 .39 5 .25 2. .37

2 23332 38464 91 .74 5. .31 2. ,97

3 23333 38464 91 .69 5, .27 3. ,05

4 23335 38464 91 .47 5. . 19 3. 36

5 23337 38464 91 .31 5. .14 3. 57

6 23339 38464 31 .10 5. ,15 3. 77

7 23340 38464 91 .09 5. ,22 3. 71

8 23343 38464 91 .33 5. ,17 3. 52

9 23344 38464 31 .54 5. 21 3. 27

10 23345 38464 91. .74 5. 27 3. 01

11 23346 38464 91, .88 5. 21 2. 92

12 23347 38464 91. .88 5. 21 2. 92

13 23348 38464 92. .01 5. 26 2. 75

14 23349 38464 92. 10 5. 41 2. 51

15 23350 38464 92. 48 5. 35 2. 18

16 23352 38464 92. 62 5. 34 2. 05

Table C-l (continued)

Point # Position,y Composition, a/o

X Y Ni A1 Ti

17 23354 38464 92 .82 5 .34 1.86

18 23356 38464 93 .15 5 .22 1.64

19 23358 38464 93 .36 5 .15 1.50

20 23362 38464 93 .55 5 .06 1.40

21 23364 38464 93 .76 4 .98 1.27

22 23366 38464 93 .66 5 .11 1.24

23 23368 38464 93 .66 5 .09 1.26

24 23370 38464 93 .81 5 .11 1.30

25 23372 38464 93 .57 5 .12 1.32

26 23374 38464 93 .40 5 .16 1.44

27 23376 38464 93 .28 5 .20 1.52

28 23378 38464 93 .08 5 .28 1.65

29 23380 38464 92 .99 5 .29 1.74

30 23382 38464 93 .08 5 .25 1.69

31 23384 38464 92 .92 5 .30 1.80

32 23386 38464 92 .73 5 .29 1.99

33 23388 38464 92 .52 5 .25 2.24

34 23390 38464 92 .06 5. .28 2.67

35 23392 38464 91 .78 .5 .25 2.98

36 23395 38464 91 .67 5 .22 3.12

37 23397 38464 91 .71 5 .18 3.13

38 23402 38464 92 .32 5 .26 2.43

39 23409 38464 92 .32 5 .34 2.35

Table C-1 (continued)

Point # Position, ]i

X Y

Composition,

Ni A1

a/o

Ti

40 23413 38464 92.51 5.27 2.24

41 23417 38464 92.64 5.34 2.03

42 23421 38464 92.58 5.33 2.10

43 23425 38464 92.81 5.28 1.92

44 23429 38464 92.96 5.20 1.85

45 23432 38464 93.32 5.16 1.53

46 23436 38464 93.57 5.13 1.31

47 23440 38464 93.69 5.11 1.21

48 23441 38464 93.77 5.03 ' 1.21

49 23442 38464 93.76 5.03 1.21

50 23444 38464 93.61 5.10 1.30

51 23448 38464 93.22 5.21 1.57

52 23452 38464 92.90 5.29 1.82

53 23456 38464 92.47 5.38 2.17

54 23460 38464 93.09 5.18 1.74

55 23462 38464 93.09 5.24 1.68

56 23464 33464 93.26 5.21 1.54

57 23466 38464 93.42 5.15 1.43

58 23468 38464 93.46 5.15 1.40

59 23470 38464 93.43 5.15 1.43

60 23472 38464 93.55 5.09 1.36

61 23476 38464 93.55 5.11 1.34

TABLE C-2

MICROSEGREGATION DATA FOR SAMPLE #1: DENDRITE B

Secondary Arm Spacing: 20y


X = 31023, Y = 32947

Point # Position, y Composition, a/o

X Y Ni A1 Ti

1 31023 32947 92.54 5.18 2.30

2 31028 32947 91.94 5.13 2.94

3 31033 32947 91.70 5.17 3.15

4 31036 32947 91.78 5.20 3.03

5 31039 32947 92.05 5.24 2.73

6 31043 32947 92.44 5.19 2.39

7 31045 32947 92.58 5.34 2.09

8 31051 32947 93.01 5.23 1.78

9 31055 32947 93.26 5.16 1.59

10 31059 32947 93.42 5.14 1.45

11 31063 32947 93.60 5.09 1.32

12 31066 32947 93.62 5.08 1.31

13 31069 32947 93.67 5.07 1.26

14 31073 32947 93.58 5.10 1.33

15 31078 32947 93.25 5.22 1.54

16 31086 32947 91.86 5.13 3.02


Point # Position, ]7 a/o

Ti

Composition,

X Y Ni A1

17 31090 32947 91. 86 4.98

18 31093 32947 91. 78 5.08

19 31097 32947 91. 87 5.12

20 31101 32947 92. 13 5.20

21 31107 32947 92. 53 5.23

22 31113 32947 93. 20 5 .11

23 31117 32947 93. 35 5.12

24 31121 32947 93. 50 5.05

25 31126 32947 93. 65 5.01

26 31134 32947 93. 83 4.94

27 31136 32947 93. 79 4.94

28 31140 32947 93. 60 5.07

29 31143 32947 93. 48 5.12

30 31145 32947 92. 72 5.09

31 31148 32947 92. 76 5.04

32 31153 32947 92. 35 4.99

3.1

3.1

3.0

2 . 6

2 . 2

1.6

1.5

1.4

1.3

1.2

1.2

1.3

1.4

2 . 2

2 . 2

2 . 6

176

TABLE C-3

MICROSEGREGATION DATA FOR SAMPLE #l:n PRIMARY DENDRITES

(v 8S* Starting Position of Point #1 *Vu ^

X = 37535, Y = 14762 7 iTt r—' •


X Y Ni A1 Ti

1 37535 14762 93.00 5.27 1.73

2 37529 14754 92.43 5.24 2.34

3 37532 14751 92.34 5.15 2.52

4 37536 14748 92.14 5.25 2.63

5 37520 14747 91.89 5.07 3.05

6 37515 14738 92.33 5.05 2.64

7 37510 14727 93.11 5 .12 1.79

8 37506 14727 93.59 4.99 1.42

9 37507 14720 93.81 4.90 1.30

10 37503 14717 93.86 4.91 1.23

11 37495 14714 93.87 4 .90 1.24

12 37497 14709 93.86 4.89 1.25

13 37485 14704 93.78 4.94 1.29

14 37488 14 708 93.74 4.94 1.32

15 37488 14704 93.57 5.00 1.44

16 37481 14697 92.97 5.21 1.83


a/o

Ti

Point # Position, \i Composition,

X Y Ni A1

17 37472 14690 92.04 5 .19

18 37475 14677 92.03 5 .02

19 37468 14673 91.95 5 .07

20 37469 14671 92.24 5 .02

21 37452 14666 92.51 5 .05

22 37448 14659 93.51 5 .01

23 37441 14657 93.72 5 .00

24 37425 14644 93.91 4 .87

25 37420 14644 93.74 4 .92

26 37416 14634 93.70 5 .01

27 37405 14624 92.58 5 .10

28 37398 14617 91.71 5 .06

29 37394 14612 91.58 5 .00

2.7

2.9

3.0

2.7

2.4

1.4

1.2

1.2

1.3

2.3

2.3

3.2

3.4

TABLE C-4

MICROSEGREGATION DATA FOR SAMPLE #2: DENDRITE A



X - 20452, Y = 38234

Point # Position, p Composition, a/o

X Y Ni A1 Ti

1 20452 38254 91.72 5.88 2.71

2 20452 38261 92.01 5.69 2.32

3 20452 38266 92.07 5.77 2.18

4 20452 38271 92.44 5.68 1.90

5 20452 38276 92.52 5.62 1.87

6 20452 38281 92.56 5.66 1.79

7 20452 38286 82.76 5.58 1.67

8 20452 38291 92.81 5.54 1.66

9 20452 38297 92.79 5.61 1.61

10 20452 38302 92.70 5.67 1.65

11 20452 38307 92.55 5.73 1.73

12 20452 38313 92.37 5.69 1.95

13 20452 38318 92.38 5.56 2.07

14 20452 38323 92.03 5.62 2.36

15 20452 38328 92.10 5.49 2.43

16 20452 38333 92.08 5.54 2.39


Point # Position, y

X Y

Composition,

Ni A1

a/o

Ni

17 20452 38338 92 .10 5.67 2.25

18 20452 38343 92 .32 5.72 1.98

19 20452 38349 92 .55 5.67 1.80

20 20452 38354 92 .71 5.62 1.68

21 20452 38359 92 .70 5.59 1.64

22 20452 38365 92 .70 5.71 1.60

23 20452 38370 92 .76 5.65 1.61

24 20452 38375 92 . 66 5.75 1.60

25 20452 38384 92 .25 5.84 1.92

26 20452 38389 92 .13 5.69 2.19

27 20452 38394 92 .19 5.42 2.41

28 20452 38400 92 .03 5.39 2.60

29 20452 38403 91 .73 5.67 2.61

30 20452 38406 92 .02 5.66 2.43

31 20452 38417 92 .26 5.72 2.03

32 20452 38423 92 .47 5.68 1.87

33 20452 38433 92 . 64 5.62 1.75

34 20452 38443 92 .74 5.62 1.65

TABLE C-5

MICROSEGREGATION DATA FOR SAMPLE #2: DENDRITE


Starting Position for Point #1:

X = 19773, Y = 39079


X Y Ni A1 Ti

1 19778 39079 92.19 5.76 2.07

2 19783 39079 92.27 5.67 2.08

3 19793 39079 91.82 5.77 2.43

4 19799 39079 91.76 5.73 2.52

5 19804 39079 91.39 5.97 2.67

6 19809 39079 91.94 5.68 2.40

7 19818 39079 92.45 5.69 1.87

8 19827 39079 92.65 5.62 1.74

9 19834 39079 92.81 5.55 1.64

10 19839 39079 92.79 5.55 1.68

11 19844 39079 92.92 5.50 1.60

12 19849 39079 92.82 5.51 1.68

13 19860 39079 92.34 5.70 1.97

14 19870 39079 92.14 5.71 2.16

15 19876 39079 91.96 5.70 2.36

16 19881 39079 91.96 5.54 2.52

25


Point # Position, y

X Y

Composition,

Ni Al

a/o

Ti

17 19887 39079 91.27 5 .84 2.91

18 19893 39079 91.88 5 .41 2.72

19 19901 39079 92.04 5 .59 2.39

20 19907 39079 92.42 5 .69 1.90

21 19913 39079 92.73 5 .58 1.70

22 19918 39079 92.79 5 .52 1.70

23 19923 39079 92.75 5 .59 1.68

24 19928 39079 92.65 5 .53 1.83

25 19933 39079 92.31 5 .70 2.01

182

TABLE C-6




X = 13206, Y = 38153

Point # Position, u Composition, a/o

X Y Ni A1

1 13206 38153 87.35 12, ,72

2 13211 38153 87.49 12. ,58

3 13217 38153 87.68 12. ,39

4 13222 38153 87.65 12. ,41

5 13229 38153 88.18 11. 88

6 13231 38153 88.35 11. 70

7 13233 38153 88.73 11. 33

8 13235 38153 88.80 11. 21

9 13237 38153 88.84 11. 21

10 13239 38153 89.06 10. 99

11 13241 38153 90.12 9. 90

12 13243 38153 90.22 9. 82

13 13245 38153 91.18 8. 85

14 13247 39153 92.02 8. 49

15 13249 38153 90.94 9. 09

16 13251 38153 90.84 9. 18

17 13253 38153 90.41 9. 63



X Y Ni Al

18 13255 38153 89.23 10.23

19 13257 38153 89.97 10.06

20 13259 38153 88.34 11.72

21 13261 38153 88.47 11.59

22 13263 38153 88.08 11.98

23 13267 38153 87.96 12.08

24 13269 38153 87.93 12.13

25 13271 38153 87.90 12.16

26 13273 38153 87.37 12.70

27 13275 38153 87.65 12.41

28 13277 38153 87.25 12.82

184

TABLE C-7




X = 22367, Y = 30696


X Y Ni A1 Ti

1 22367 30696 87. 98 11.63 .45

2 22367 30702 87. 87 11.73 .45

3 22367 30706 87. 88 11.72 .47

4 22367 30710 87. 87 11.66 .52

5 22367 30715 87. 85 11.73 .48

6 22367 30719 87. 90 11.69 .46

7 22367 30722 88. 14 11.47 .44

8 22367 30726 88. 22 11.38 .46

9 22367 30730 88. 68 11.03 .34

10 22367 30749 89. 26 10.46 .32

11 22367 30753 89. 10 10.47 .28

12 22367 30758 89. 18 10.59 .28

13 22367 30765 89. 68 10.14 .22

14 22367 30768 89. 89 9.92 .23

15 22367 30775 90. 26 9.57 .20

16 22367 30778 90. 00 9.85 .19

17 22367 30780 91. 74 8.15 .13

24 3o



Ti X Y Ni A1

18 23367 30783 91.07 8. 79

19 23367 30787 90.43 9. 46

20 23367 30791 89.95 9. 88

21 22367 30806 89.02 10. 78

22 22367 30812 89.50 10. 31

23 22367 30818 89.85 10. 03

24 22367 30824 88.90 10. 82

25 22367 30842 88.57 11. 09

26 22367 30845 88.43 11. 19

27 22367 30848 88.33 11. 31

28 22367 30851 88.38 11. 24

29 22367 30854 88.36 11. 23

30 22367 30858 88.48 11. 17

31 22367 30862 88.87 10. 83

. 1

. 1

. 2

.2

. 2

. 1

.3

.3

.4

.4

.4

.4

.4

.3

186

TABLE C-8


Secondary Arm Spacing: 33p

Starting Position for Point #1

X = 23019, Y = 30254


X Y Ni A1 Ti

1 23019 30254 87.87 10.37 1.81

2 23025 30254 87. 71 10.50 1.85

3 23031 30254 87.66 10.54 1.86

4 23036 30254 87.68 10.47 1.91

5 23042 30254 87.63 10.48 1.95

6 23045 30254 87.70 10.39 1.97

7 23050 30254 87.72 10.41 1.93

8 23054 30254 87.97 10.31 1.77

9 23065 30254 89.28 9.54 1.23

10 23074 30254 89.67 9.34 1.03

11 23077 30254 89.62 9.42 1.00

12 23087 30254 89.79 9.27 0.99

13 23093 30254 89.78 9.28 0.98

14 23099 30254 89.75 9.33 0.96

15 23105 30254 90.17 9.94 0.92

16 23110 30254 89.51 9.55 0.97



X Y Ni A1 Ti

17 23117 30254 89 .41 9.62 1.02

18 23125 30254 89 .28 9.72 1.04

19 23131• 30254 89 .38 9.70 1.02

20 23138 30254 89 .08 9.86 1.11

21 23143 30254 89 .49 9.52 1.03

22 23151 30254 89 .25 9.63 1.16

23 23170 30271 89 .06 9.67 1.32

24 23177 30271 87 .68 10.58 1.08

TABLE C-9



Starting Position for Point 1:

X = 22633, Y = 30701

Point # Position, ]i Composition, a/o

X Y Ni A1 Ti

1 22633 30701 89,25 9.70 1.10

2 22654 30701 89.92 9.20 0.92

3 22669 30701 89. 76 9.43 0.85

4 22682 30701 90.40 8.99 0.65

5 22689 30701 90.44 8.96 0.63

6 22691 30701 90.42 9.00 0.62

7 22698 30701 90.53 8.91 0.58

8 22705 30701 90.35 9.05 0.63

9 22722 30701 90.42 9.00 0.62

10 22733 30701 90.61 8.80 0.62

11 22752 30701 90.11 9.19 0.73

12 22773 30701 89.83 9.37 0.84

13 22791 30701 89.94 9.23 0.87

14 22800 30701 88.99 9.97 1.08

TABLE C-10



Starting Position for Point

X = 15482, Y = 30900 '

#1:

I 5V


X Y Ni A1 Ti

1 15482 30900 87.85 9.82 2.26

2 15482 30895 87.95 9.80 2.20

3 15482 30890 87.86 9.87 2.26

4 15482 30885 87.71 10.00 2.25

5 15482 30880 87.87 9.86 2.24

6 15482 30875 88 .03 9.79 2.18

7 15482 30870 88 .60 9.29 2.06

8 15482 30865 88.44 9.59 1.93

9 15482 30860 88.81 9.29 1.85

10 15482 30855 89.04 9.28 1.65

11 15482 30850 90.02 8.47 1.44

12 15482 30845 90.59 7.99 1.37

13 15482 30840 89.93 8.34 1.69

14 15482 30835 89.79 8.52 1.62

15 15482 30830 89.74 8.62 1.61

16 15482 30825 88.89 9.46 1.61

Table C-10(continued)


X Y Ni A1 Ti

17 15482 30820 88.72 9.59 1.66

18 15482 30815 88.59 9.59 1.82

19 15482 30810 87.78 9.55 2.61

20 15482 30805 87.63 9.81 2.49

191

TABLE C-ll



Starting Position for Point

X = 20320, Y = 31682

# 1:


X •Y Ni . A1 Ti

1 20320 31682 87 .91 9 .76 2.38

2 20325 31682 87 .55 9 .85 2.67

3 20330 31682 87 .30 10 .02 2.73

4 20335 31682 87 .13 10 .03 2.91

5 20340 31682 87 .39 9 .99 2.68

6 20345 31682 87 .29 10 .08 2.69

7 20350 31682 87 .88 9 .82 2.36

8 20355 31682 89 .29 9 .18 1.58

9 20360 31682 89 .12 9 .09 1.84

10 20365 31682 89 .05 9 .25 1.75

11 20370 31682 90 .01 8 .48 1.54

12 20375 31682 88 .78 9 .43 1.84

13 20380 31682 88 .74 9 .49 1.82

14 20385 31682 88 .30 9 .75 2.00

15 20390 31682 88 .29 9 .68 2.08

16 20395 31682 87 .98 9 .95 2.13

17 20400 31682 87 .96 9 .94 2.15

192

TABLE C-12




X = 19208, Y = 35204


X Y Ni A1 Ti

1 19208 35213 87.59 5.17 7.27

2 19208 35214 87.56 5.08 7.39

3 19208 35215 87.47 5.09 7.49

4 19208 35216 87.07 5.19 7.78

5 19208 35217 87.20 5.15 7.70

6 19208 35218 87.78 5.23 8.03

7 19208 35220 86.64 5.29 8.40

8 19208 35222 86.26 5.39 8.40

9 19208 35223 86.72 5.22 8.10

10 19208 35224 86.73 5.29 8.02

11 19208 35228 86.74 5.33 7.97

12 19208 35232 87.09 5.45 7.51

13 19208 35242 89.72 5.13 5.18

14 19208 35244 91.55 4.43 4.04

15 19208 35251 91.44 4.67 3.91

16 19208 35261 90 .73 5.09 4,20

'34. IZ



X Y Ni A1 Ti

17 19208 35266 90 .44 5. 25 4 .34

18 19208 35271 90 .22 5. 26 4 .54

19 19208 35276 90 .08 5. 40 4 .54

20 19208 35282 91 .19 4. 90 3 .92

21 19208 35285 90 .43 5. 01 4 .59

22 19208 35294 87 .41 5. 62 7 .01

23 19208 35296 87 .47 5. 56 7 .01

24 19208 35299 88 .86 5. 00 6 .17

25 19208 35302 88 .40 5. 26 6 .37

26 19208 35304 88 .73 + 96 6 .34

27 19208 35310 88 .94 4. 81 6 .28

28 19208 35312 87 .90 5. 29 6 .85

29 19208 35325 88 .48 5. 22 6 .33

30 19208 35339 90 .49 5. 38 4 .15

31 19208 35356 90 .13 5. 65 4 .26

32 19208 35361 91 .34 5. 11 3 .75

33 19208 35374 91 .63 4. 64 3 .75

34 19208 35377 90 .43 , 4. 98 4 .61

35 19208 35384 90 .42 4. 86 4 .74

36 19208 35389 90 .02 4. 92 5 .08

194

TABLE C-13




X = 7904, Y = 41544


X Y Ni A1 Ti

1 7904 41544 87.72 5.37 6.95

2 7904 41548 86.78 5.21 7.85

3 7904 41556 85.72 5.32 8.01

4 7904 41561 87.18 5.19 7.67

5 7904 41576 87.58 5.34 7.12

6 7904 41578 89.35 5 . 0 2 5.66

7 7904 41580 90.93 4.73 4.35

8 7904 41584 90.98 5.17 3.88

9 7904 41592 91.06 5.20 3.76

10 7904 41603 90.20 5.40 4.43

11 7904 41606 90.25 5.08 4.69

12 7904 41614 87 .58 5.15 7.31

13 7904 41622 87.59 5.03 7.42

14 7904 41630 87.87 5.11 7.05

15 7904 41638 88.77 5.02 6.25

16 7904 41640 89.45 5.12 5.45



X Y Ni A1 Ti

17 7904 41645 90.93 4.73 4.35

18 7904 41650 90.98 5.17 3.88

19 7904 41655 91.06 5.20 3.76

20 7904 41660 90.25 5.08 4.69

TABLE C-14


Secondary Arm Spacing; 18y


X = 25683, Y = 39910


X Y Ni A1 Ti

1 25683 39910 89 .33 2.82 7.86

2 25683 39913 89 .06 2.73 8.22

. 3 25683 39916 89 .03 2.60 8.38

4 25683 39919 88.31 2.66 9.04

5 25683 39928 88.49 2.76 8.76

6 25683 39933 90.85 2.49 6.66

7 25683 39938 90.54 2.83 6.64

8 25683 39943 90.86 2.88 6.27

9 25683 39948 90.51 2.95 6.55

10 25683 39953 90.48 2.94 6.59

11 25683 39958 90.50 2.85 6.66

12 25683 39960 89.52 2.96 7.53

13 25683 39964 88.70 2.98 8.34

14 25683 39968 88.91 2.42 8.68

15 25683 39972 88.59 2.41 9.00

16 25683 39976 89.08 2.41 8.52



X Y Ni A1 Ti

17 25683 39978 89.54 2.79 7.68

18 25683 39982 90.63 2.93 6.44

19 25683 39987 90.62 2.90 6.49

20 25683 39992 91.09 2.66 6.26

21 25683 39998 90.45 2.78 6.77

22 25683 40010 90.28 2.49 7.24

23 25683 40013 89.92 2.42 7.67

24 '25683 40015 89 .08 2.37 8.56

25 25683 40019 88:87 2.54 8.60

26 25683 40022 89.74 2.74 7.53

27 25683 40027 89.72 2.97 7.33

28 25683 40032 90.55 2.90 6.57

29 25683 40037 90.85 2.92 6.24

30 25683 40042 90.84 2.85 6.33

31 25683 40047 90.84 2.87 6.30

TABLE C-15




X =12524, Y = 40808


X Y Ni Al Ti

1 12524 40808 90.99 2.59 6.43

2 12524 40816 91.21 2.55 6.25

3 12524 40820 91.29 2.59 6.13

4 12524 40823 91.17 2.58 . 6.25

5 12524 40829 90.80 2.61 6.60

6 12524 40836 90.22 2.29 7.49

7 12524 40841 89.97 2.11 7.92

8 12524 40846 89.65 2.20 8.15

9 12524 40851 89.79 2.24 7.98

10 12524 40856 90.45 2.26 7 .30

11 12524 40862 90.51 2.45 7.05

12 12524 40872 91.16 2.57 6.28

13 12524 40880 81.15 2.59 6.26

14 12524 40887 81.21 2.60 6.19

15 12524 40897 91.07 2.58 6.36

16 12524 40906 90.40 2.33 7.28



X Y Ni A1 Ti

17 12524 40918 89.38 2.22 8 .41

18 12524 40923 89.66 2.20 8.15

19 12524 40928 90.15 2.28 7.57

20 12524 40936 90.40 2.44 7.17

200

TABLE C-16


Secondary Arm Spacing: 33u


X = 13275, Y = 42245

Point # Position, \i Composition, a/o

X Y Ni A1 Ti

1 13275 42245 89.67 2.15 8.19

2 13275 42259 90.37 2.37 7.27

3 13275 42264 90.85 2.39 6.76

4 13275 42269 91.11 2.47 6.43

5 13275 42274 91.16 2.55 6.29

6 13275 42279 91.12 2.53 6.35

7 13275 42284 90.77 2.43 6.81

8 13275 42292 90.36 2.25 7.40

9 13275 42300 89.87 2.22 7.92

10 13275 42308 90.20 2.14 7.66

11 13275 4231S 90.20 2.36 7.45

12 13275 42323 90.60 2.47 6.94

13 13275 42328 91.07 2.55 6.39

14 13275 42333 91.30 2.54 6.16

15 13275 42338 91.20 2.58 6.23

16 13275 42343 90.91 2.49 6.61



X Y Ni A1 Ti

17 13275 42350 89.94 2.28 7.79

18 13275 42357 90.23 2.18 7.59

19 13275 42364 90.19 2.32 7.49

20 13275 42371 90.55 2.37 7.08

21 13275 42376 91.12 2.49 6.40

22 13275 42381 81.28 2.51 6.22

23 13275 42386 91.32 2.54 6.14

24 13275 42390 91.33 2.59 6.08

25 13275 _ 42401 90.81 2.46 6.74

26 13275 42407 90.26 2.21 7,53

27 13275 42414 90.49 2.18 7.34

28 13275 42421 90.53 2.25 7.23

TABLE C-17


Secondary Arm Spacing: 38u •> V—' I—J

Starting Position of Point #1: v f—\ t—\

X = 10637, Y = 32052 W &

Poin t # P o s i t i o n , y C o m p o s i t i o n , a/ o

X Y Ni Ti

1 10637 32052 89.60 10.38

2 10637 32055 89.61 10.37

3 10637 30259 89.68 10.31

4 10637 30267 90.25 9.73

5 10637 32077 90.89 9.10

6 10637 32084 91.0- 8.90

7 10637 32089 91.12 8.87

8 10637 32096 91.51 8.47

9 10637 32102 91.59 8.40

10 10637 32105 91.65 8.34

11 10637 32108 91.56 8.42

12 10637 32112 81.58 8.41

13 10637 32118 91.52 8.48

14 10637 32127 91.50 8.49

15 10637 32133 91.41 8.58

16 10637 32139 91.07 8.91


Point # Position, y

X Y

17 10637 32144

18 10637 32151

19 10637 32160

20 10637 32164

21 10637 32167

22 10637 32171

23 10637 32176

24 10637 32181

25 10637 32186

26 10637 32191

27 10637 32199

28 10637 32206

29 10637 32213

30 10637 32219

31 10637 32225

32 10637 32235

33 10637 32241

34 10637 32248

35 10637 32254

36 10637 32263

37 10637 32273

38 10637 32279

39 10637 32285

Composition, a/o

Ni Ti

90.97 9.02

89.76 10.22

89.64 10.34

87.34 12.64

87.18 12.80

87.65 12.34

87.31 12.67

87.76 12.22

87.23 12.76

88.75 11.24

89.90 10.09

90.44 9.54

90.98 9.00

91.26 8.72

91.48 8.51

91.54 8.45

91.50 8.49

91.66 8.33

91.62 8.37

91.27 8.71

90.96 9.02

90.25 9.73

89.42 10.57



X Y Ni Ti

40 10637 32291 88.82 11.16

41 10637 32297 88.50 11.48

42 10637 32303 88.61 11.37

43 10637 32307 88.45 11.54

44 10637 32311 88.05 11.93

45 10637 32315 87.39 12.59

46 10637 32321 87.75 12.23

205

TABLE C-18


Secondary Arm Spacing: 38y —J >—' *—' '—' \

Starting Position of Point #1: \ <—\ | \ ?

X = 16716, Y = 31459 w ^


X Y Ni Ti

1 16716 31459 87.75 12.23

2 16716 31451 87 .98 12.02

3 16716 31445 87.94 12.04

4 16716 31440 88.21 11.78

5 16716 31435 88.18 11.81

6 16716 31428 88.38 11.61

7 16716 31418 88.70 11.29

8 16716 31409 88.90 11.09

9 16716 31400 89.54 10.45

10 16716 31392 89.90 10.09

11 16716 31382 90.34 9.65

12 16716 31377 90.33 9.66

13 16716 31372 90.45 9.54

14 16716 31367 90.48 9.50

15 16716 31365 90.42 9 .56

16 16716 31359 90.44 9.55



X Y Ni Ti

17 16716 31351 90.22 9.76

18 16716 31343 89.89 10.09

19 16716 31335 89.64 10.34

20 16716 31322 88.41 11.59

21 16716 31309 87.58 12.40

TABLE C-19


Secondary Arm Spacing: 19p


X = 16009, Y = 23034


X Y Ni Ti

1 16009 23048 87.80 12.28

2 16009 23053 87.75 12.24

3 16009 23057 87.52 12.47

4 16009 23062 87.62 12.37

5 16009 23065 87.82 12.16

6 16009 23067 87.99 11.97

7 16009 23072 88.28 11. 70

8 16009 23079 89.50 10.49

9 16009 23087 90.11 9.87

10 16009 23092 90.81 9.18

11 16009 23100 91.27 8.71

12 16009 23106 91.42 8.57

13 16009 23110 91.33 8.65

14 16009 23115 91.27 8.72

15 16009 23120 90.82 9.16

16 16009 23129 98.66 10.33

17 16009 23136 88.14 11.85

lo



X Y Ni Ti

18 16009 23144 83.25 12. 76

19 16009 23151 87.12 12.87

20 16009 23156 87.69 12.29

209

TABLE C-20




X = 15654, Y = 26249

Point # Position, n Composition, a/o

X Y Ni Ti

1 15654 26286 83.42 12.57

2 15654 26281 87.54 12.44

3 15654 26276 87.68 12.31

4 15654 26271 87.62 12.36

5 15654 26267 87.98 12.01

6 15654 26266 88.84 11.15

7 15654 26265 90.74 9.24

8 15654 26264 90.87 9.12

9 15654 26263 90.95 9.03

10 15654 26262 91.26 8.72

11 15654 26258 91.34 8.65

12 15654 26254 91.36 8.63

13 15654 26250 91.34 8.65

14 15654 26246 91.19 8.79

15 15654 26241 91.02 8.94

16 15654 26236 90.88 9.10

17 15654 26232 90.78 9.21



X Y Ni Ti

18 15654 26227 90.46 9.44

19 15654 26223 90.28 9.71

20 15654 26216 89.88 10.11

21 15654 26213 89.54 10.44

22 15654 26208 88.96 11.02

23 15654 26205 88.65 11.33

24 15654 26202 88.53 11.36

25 15654 26199 88.54 11.44

26 15654 26196 88.52 11.47

2

3

4

5

6

7

8

9

10

11

12

13

14

15

16

17

LIST OF REFERENCES

W. C. Bigelow: Am. Soc. Met. Spec. Tech. Publ., 1958, vol. 245, p.73.

J. R. Mihalisin and D. F. Decker: Trans. TMS-AIME, 1960, vol. 218, p.507.

R. L. Stegman, P. Shahiman, and M. R. Achter: Trans. AIME, 1969, vol. 245, p.1759.

M. Volmer and A. Weber: Z. Phys. Chem. , 1926, vol. 119, p.227.

R. Becker and W. Doring: Ann. Phys., 1935, vol. 24, p.719.

D. Turnbull: J. Metals, 1950, vol. 188, p.1144.

J. L. Walker: Liquid Metals and Solidification, Amer. Soc. Metals, Metals Park, Ohio, 1958.

D. Turnbull: J. Chem. Phys., 1950, vol. 18, p.198.

M. Volmer: Z. Electrochem., 1929, vol. 35, p.555.

B. E. Sundquist and L. F. Mondolfo: Trans. AIME, 1961, vo1. 221, p.15 7.

G. H. Gulliver: Metallic Alloys, C. Griffin and Co., London, 1922.

E. Scheil: Z. Metallk, 1942, vol. 34, p.70.

W. G. Pfann: Trans. AIME, 1952, vol. 194, p.747.

M. C. Flemings: Solidification Processing, McGraw-Hill, New York, 1974.

N. J. Petch: J.'Iron Steel Inst., 1953, vol. 173, p.25.

J. F. Wallace: Elec. Conf. Proc. AIME, 1962, vol. 20. p.125.

E. M. Passmore, M. C. Flemings, and H. F. Taylor: Trans. AFS, 1958, vol. 66, p.96.

211

18)

19)

20)

21)

22)

23)

24)

25)

26)

27)

28)

29)

30)

31)

32)

33)

212

S. F. Frederick and W. A. Baily: Trans. AIME, 1968, vol. 242, p.2063.

S. N. Singh, B. P. Bardes, and M. C. Flemings: Met. Trans., 1970, vol. 1, p.1383.

F. Weinberg and R. K. Buhr: The Solidification of Metals, ISI Publ. 110, 1968, p.295.

R. D. Doherty and D. H. Melford: J. Iron Steel Inst., 1966, vol. 204, p.1131.

M. C. Flemings: The Solidification of Metals, ISI Publ. 110, 1968, p.T04~

P. J. Ahearne and F. C. Quigley: J. Iron Steel Inst., 1966, vol. 204, p.16.

A. Kohn: Rev. Met., 1970, vol. 67, p.23.

L. S. Darken and R. W. Gurry: Physical Chemistry of Metals, McGraw-Hill, New York, 1953.

R. Hultgren, P. D. Desai, D. T. Hawkins, M. Gleiser, and K. K. Kelly: Selected Values of the Thermodynamic Properties of Binary Alloys, Amer. Soc. Metals, Metals Park, Ohio, 1973.

M. Hansen: Constitution of Binary Alloys, 2nd ed., McGraw-Hill, New York, 1958.

R. P. Elliott: Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York, 1968.

F. A. Shunk: Constitution of Binary Alloys, Third Supplement, McGraw-Hill, New York, 1969.

G. K. Sigworth, J. F. Elliott, G. A. Vaughn, and G. H. Geiger: Can. Met. Quart., in press.

E. T. Turkdogan: J. Iron Steel Inst., 1972, vol. 210. p.21.

K. Torsell and M. Olette: Comptes Rendus, 1969, vol. 268, p.81.

J. F. Elliott, Y. Iguchi, and T. H. Chiang: Proc. Elec. Furnace Steel Making Conf., 1977, vol. 37, p . 2 0 0 .

34)

35)

36)

37)

38)

39)

40)

41)

42)

43)

44)

45)

46)

47)

213

K. R. Olen, L. S. Gonano, and E. L. Heck, J. Metals, 1970, vol. 8, p.36.

K. Oberg, S. Eketorp, I. Isaksson, C. Nilsson, and A. S. Jain: Jernkont. Ann., 1971, vol. 155, p.17.

E. Snape and P. R. Beeley: Modern Casting, 1969, vol. 6, p.159.

J. T. Wallace: Solidification Technology, J. L. Burke, M. C. Flemings, and A. E. Gorum, eds., Brook-Hill, Chestnut Hill, Massachusetts, 1974, p . 15 .

H. N. Cummings, F. B. Stulen, and W. C. Shulte: ASM Trans., 1957, vol. 49, p.482.

A. Taylor and R. W. Floyd: J. Inst. Metals, 1952-53, vol. 81, p.451.

J. Gurland, Quantitative Microscopy, pp.344-352, McGraw-Hill, New York, 1968.

0. D. Sherby and M. T. Simnad: Trans. ASM, 1961, vol. 54, p.227.

A. K. Bhambri, T. Z. Kattamis, J. E. Morral: Met. Trans., 1978, vol. 6B, p.523.

T. F. Bower, H. D. Brody, and M. C. Flemings: Trans. AIME, 1966, vol. 236, p.624.

W. E. Brower and M. C. Flemings: Solidification of Iron Base Alloys at Large Degrees o± Undercooling, AMMRC CR-69-14/1, Contract No. DA-46-68-C-0044, Watertown, Massachusetts, 1969.

H. D. Brody: Solidification Technology, J. L. Burke, M. C. Flemings, and A. E. Gorum, eds., Brook-Hill, Chestnut Hill, Massachusetts, 1974, p.53.

T. B. Braun, J. F. Elliott, M. C. Flemings: Trans. Met., in press.

J. W. Colby: Magic IV - A Computer Program for Quantitative Electron Microprobe Analysis, Bell Telephone Lab., Inc., Allentown, Pennsylvania.

214

48) P. Duncumb and S. J. B. Reed: Quantitative Electron Probe Microanalysis, NBS Special Publ. 298, K. F. J. Iieinrich, eds. , 1968.

49) D. R. Beaman and J. A. Isasi: Electron Beam Microanalysis, ASTM Special Tech. Publ. 506, Philadelphia, Pennsylvania, 1974.

50) W. Bambynek: Rev. Mod. Phys., 1972, vol. 4, p.716.

solidification of nickel-base alloys containing … · 2020. 4. 2. · solidification of...

Documents