NBSIR 73-415

Noble Metal Constitution Diagrams:

Part II

R. M. Waterstrat and R. C. Manuszewski

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

Final Report September 1973

Issued August 1975

Prepared for

American Dental Association Research Unit

at the National Bureau of Standards,

Washington, D. C. 20234

NBS-114A (REV. 7-73)

U .S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

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1. PUBLICATION OR REPORT NO.

NBSIR 73-415

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4. TITLE AND SUBTITLE

Noble Metal Constitution Diagrams: Part II

5. Publication Date

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7. AUTHOR(S)

R. M. Waterstrat and R. C. Manuszewski

8. Performing Organ. Report No.

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NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 3119585

11« Contract /Grant No. DE02455

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American Dental Association

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Chicago, Illinois 60611

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15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual summary of most si^ificant information. If document includes a significant ' bibliography or literature survey, mention it here.)

Six binary constitution diagrams involving the noble metals

are presented. These diagrams include the V-Ru, V-Rh, Nb-Os,

Nb-Pd, Nb-Pt and Ta-Pd alloy systems.

Experimental alloys were prepared from starting materials having

a nominal purity of at least 99.9% and precautions weire taken to

insure that no significant contamination was introduced during

alloy preparation and heat-treatment. Temperatures were

measured to an accuracy within ±20°C.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons)

Alloy phases; constitution diagrams; phase diagrams; phase equilibria;

platinum group metals ; noble metals

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NBSIR 73-415

NOBLE METAL CONSTITUTION DIAGRAMS:

PART II

R. M. Waterstrat and R. C. Manuszewski

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

Final Report September 1973

Issued August 1975

Prepared for American Dental Association Research Unit

at the National Bureau of Standards,

Washington, D. C. 20234

U.S. DEPARTMENT OF COMMERCE, Rogers C.B. Morton, Secretary

James A. Baker, III. Under Secretary

Dr. Betsy Ancker-Johnson, Assistant Secretary for Science arid Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Acting Director

SEPTEMBER 28, 1973

NOBLE METAL CONSTITUTION DIAGRAMS: PART

by

R. M. Waterstrat and R. C. Manuszewski

This investigation is supported by Grant DE02455 to the American

Dental Association by the National Institute of Dental Research and

is part of the dental research program conducted by the National

Bureau of Standards in cooperation with the American Dental

Association, the National Institute of Dental Research and the

Dental Research Division of the U. S. Army Medical Research

and Development Command.

+ Noble Metal Constitution Diagrams, Part I, was issued

April 20, 1971 as NBS Report 10571.

* Research Associates, American Dental Association Research

Program in the Dental Research Section, National Bureau of

Standards, Washington, D. C. 20234.

i
NBS PROGRESS REPORT 73-415

NOBLE METAL CONSTITUTION DIAGRAMS : PART II

R. M. Waterstrat and R. C. Manuszewski

SEPTEMBER 1973

Performed under

NIDR Research Grant DE02455 to

the American Dental Association.

Conducted at the

National Bureau of Standards

Washington, D. C. 20234

ii
FOREWORD

This investigation was supported primarily by Research

Grant DE02455 to the American Dental Association from the

National Institute of Dental Research. The research was conducted at the National Bureau of Standards during the period

from September 1, 1970 to August 31, 1973.

The electron probe microanalyses of two-phase alloys was

performed at the National Bureau of Standards by Mrs. Martha

Darr, Mr. Robert Myklebust and Mr. Charles Fiore under the

direction of Dr. Kurt Heinrich.

Constitution diagrams for the binary Nb-Pd, Nb-Pt and

Ta-Pd systems are partially based on unpublis hed experimental

data which was generously made available to us by Prof. B. C.

Giessen of Northeastern University, Boston, Mass.

Earlier work in this research program , conducted under

the same Research Grant, is reported in NBS Report 10571 dated

April 20 , 1971.

iii
ABSTRACT

Six binary constitution diagrams involving the noble

metals are presented. These diagrams include the V-Ru,

V-Rh, Nb-Os, Nb-Pd, Nb- Pt and Ta- Pd alloy systems.

Experimental alloys were prepared from starting materials

having a nominal purity of at least 99. 9% and precautions were

taken to insure that no significant contamination was introduced

during alloy preparation and heat- treatment. Temperatures

were measured to an accuracy within ±20° C.

iv
TABLE OF CONTENTS

Page No.

I. SUMMARY 1 A. Constitution Diagrams 1 and 2

B. Experimental Method 3

II. INTRODUCTION 5

A. Background 5

B. Experimental Approach 6

m. VANADIUM-RUTHENIUM CONSTITUTION DIAGRAM .... 10

A. Previous Studies 10

B. Materials 10

C. Alloy Preparation 11 D. Temperature Measurement U

E. Thermal Treatments 13

F. Methods for Determination of Phase Boundaries 14

G. Experimental Results 16

H. Discussion 24

REFERENCES

rV. VANADIUM-RHODIUM CONSTITUTION DIAGRAM 26

A. Previous Studies 26

B. Materials 26

C. Alloy Preparation 26

D. Temperature Measurement 28

E. Thermal Treatments 28

F. Methods for Determination of Phase Boundaries 30

G. Experimental Results 31

H. Discussion 44

REFERENCES

.

. 51

V
TABLE OF CONTENTS (Cont'd)

Page No.

V. NIOBIUM-OS MIUM CONS TITUTION DUGRAM 53

A. Previous Studies 53

B. Materials 53

C. Alloy Preparation 55

D. Temperature Measurement 56

E. Thermal Treatments 57

F. Methods for Determination of Phase Boundaries ... 59

G. Experimental Results 62

H. Discussion 73

REFERENCES 77

VI. NIOBIUM- PALLADIUM CONSTITUTION DIAGRAM . . . 79

A. Previous Studies 79

B. Materials 80

C. Alloy Preparation 81

D. Temperature Measurement 81

E. Thermal Treatments 82

F. Methods for Determination of Phase Boundaries ... 84

G. Experimental Results 85

H. Discussion 97

REFERENCES 102

vi
TABLE OF CONTENTS (Cont'd)

Page No.

Vn. NIOBIUM- PLATINUM CONSTITUTION DIAGRAM . ...... 104

A. Previous Studies 104

B. Materials 105

C. Alloy Preparation 105

D. Temperature Measurement 106

E. Thermal Treatments 106

F. Methods for Determination of Phase Boundaries 107

G. Experimental Results 109

H. Discussion 118

REFERENCES 122

Vm. TANTALUM- PALLADIUM CONSTITUTION DIAGRAM 123

A. Previous Studies 123

B. Materials 124

C. Alloy Preparation 124

D. Temperature Measurement 125

E. Thermal Treatments 125

F. Methods for Determination of Phase Boundaries 127

G. Experimental Results 128

H. Discussion 139

REFERENCES 143

APPENDIX I. METALLOGRAPHIC PROCEDURES 144

APPENDIX II. X-RAY DIFFRACTION PROCEDURES 146

APPENDIX III. ELECTRON MICROPROBE ANALYSIS 148

APPENDIX IV. ATOMIC VOLUME RELATIONSHIPS 150

vii
LIST OF ILLUSTRATIONS

Figure

1. Constitution diagrams of V-Ru, V-Rh, Nb-Os and

2. Constitution diagrams of Nb-Pt and Ta-Pd . .

3. Vanadium- Ruthenium constitution diagram . .

4. Microstructures of vanadium- ruthenium alloys

5. Microstructures of vanadium- ruthenium alloys

6. Microstructures of vanadium- ruthenium alloys

7. Vanadium- Rhodium constitution diagram . .

8. Microstructures of vanadium- rhodium alloys.

9. Microstructures of vanadium- rhodium alloys.

10. Microstructures of vanadium- rhodium alloys

11. Microstructures of vanadium- rhodium alloys.

12. Microstructures of vanadium- rhodium alloys .

13. Microstructures of vanadium- rhodium alloys .

14. Niobium-Osmium constitution diagram

15. Microstructures of niobium-osmium alloys

16. Microstructures of niobium-osmium alloys .

17. Microstructures of niobium-osmium alloys .

Page No.

1

2

17

19

22

23

32

35

36

37

39

40

42

61

68

71

72

viii

n f
LIST OF ILLUSTRATIONS

Figure Page No.

18. Niobium- Palladium constitution diagram 86

19. Microstructures of niobium-palladium alloys ... 88

20. Microstructures of niobium-palladium alloys ... 89

21. Mic rostructures of niobium-palladium alloys ... 90

22. Microstructures of niobium-palladium alloys ... 91

23. Microstructures of niobium- palladium alloys ... 95

24. Microstructures of niobium-palladium alloys ... 96

25. Niobium- Platinum constitution diagram 110

26. Microstructures of niobium- platinum alloys ... 112 •

27. Microstructures of niobium- platinum alloys ... 113

28. Microstructures of niobium- platinum alloys ... 116

29. Microstructures of niobium- platinum alloys ... 117

30. Tantalum -Palladium constitution diagram .... 129

31. Microstructures of tantalum -palladium alloys . . 131

32. Microstructures of tantalum-palladium alloys . . 133

33. Microstructures of tantalum -palladium alloys . . 134

34. Microstructures of tantalum- palladium alloys . . 135

35. Microstructures of tantalum- palladium alloys . . 137

36. Microstructure of a tantalum -palladium alloy . . 138

ix
LIST OF ILLUSTRATIONS

Page No.

37. Atomic volumes of V-Ru alloys 151

38. Atomic volumes of V-Rh alloys 152

39. Atomic volumes of Nb-Os alloys 153

40. Atomic volumes of Nb-Pd alloys 154

41. Atomic volumes of Nb-Pt alloys 155

42. Atomic volumes of Ta-Pd alloys 156

X
LIST OF TABLES

TABLE Page No.

1. Typical Analyses of the Materials Used in

Preparir^ the Experimental Alloys 12

2. Summary of Equilibration Treatments for

Vanadium- Ruthenium Alloys 14

3. Lattice Parameters of V-Ru Alloys 20

4. Typical Analyses of the Materials Used in

Preparing the Experimental Alloys 27

5. Schedule of Typical Heat Treatments for

Vanadium- Rhodium 29

6. Relative Intensities (I) and Interplaner Spacings

(d) for a 3 Phase Lines Compared with

Equivalent Values Reported for the Phase

a VIr (5) 46 & 47

7. Invariant Reactions in the Vanadium- Rhodium

System 48

8. Crystallographic Data for Vanadium- Rhodium

Intermediate Phases 49

9. Lattice Parameters of Intermediate Phases and

Terminal Solid Solutions for Vanadium- Rhodium . 50

10. Typical Analyses of the Materials Used in

Preparing the Experimental Alloys 54

11. Summary of Typical Equilibration Treatments

for Niobium-Osmium Alloys 58

12. Crystallographic Data for Niobium-Osmium

Intermediate Phases 63

xi
LIST OF TABLES

TABLE Page No.

13. Lattice Parameters of Intermediate Phases and

Terminal Solid Solutions for Niobium-Osmium. . . 64

14. Observed and Calculated Relative Intensities

for NbggOSg^ x-P^ase 65,66 &67

15. Comparison of Site Occupancies for the y^, a and

p Phases 75

16. Typical Equilibration Treatments for NiobiumPalladium Alloys 83

17. Crystallographic Data for Niobium- Palladium

Intermediate Phases 100

18. Lattice Parameters of Intermediate Phases and

Terminal Solid Solutions for Niobium- Palladium . . 101

19. Typical Equilibration Treatments for NiobiumPlatinum Alloys 108

20. Crystallographic Data for Niobium- Platinum

Intermediate Phases 120

21. Lattice Parameters of Intermediate Phases and

Terminal Solid Solutions for Niobium- Platinum . . 121

22. Typical Equilibration Treatments for TantalumPalladium Alloys 126

23. Crystallographic Data for Tantalum -Palladium

Intermediate Phases 141

24. Lattice Parameters of Intermediate Phases and

Terminal Solid Solutions for Tantalum -Palladium . 142

xii
I. SUMMARY

A. Constitution Diagrams

The constitution diagrams which have been determined are

presented in Figures 1 and 2.

1

1 II' 1

TWO PHASE

UElT'NC bkam cm rcxriNC

MELTING CaWXra) OM HCXTWG _

PHASE 8

€ L 1 -3

/ l790Hp»c\"*"

/H

•

C Hi f

II 1 10 20 30 40 SO SO 7D BO 90 KX)

ATOMIC PERCENT \MAD<UM V

WDOfT PERCENT uUiAOuU

Fig. 1 - Constitution diagrams of vanadium- ruthenium; vanadiumrhodium; niobium- osmium; and niobium-palladium.

1
0 10 20 30 40 50 60 70 80 90 100

Pr ATOMIC PERCENT NIOBIUM Nb

Fig. 2 - Constitution diagrams of niobium- platinum and

tantalum-palladium.

2
B. Experimental Method

1. Materials

The purity of all the starting materials used in

this study was at least 99.9%. This purity was considered to be

adequate to insure a reliable constitution diagram within the

accuracy limits imposed by uncertainties in temperature and

composition measurements.

2. Alloy Preparation

The experimental alloys were prepared by nonconsumable arc- melting of the pure metals. The pure metals were

submitted to melting in various forms. Vanadium was in the form

of electrolytic dendrites which were compacted into 1/2- inch

diameter cylinders of the appropriate weight. Mixed powders of

niobium, tantalum, ruthenium, osmium and rhodium were compressed into 1/2- inch diameter cylinders. Platinum and palladium

were in the form of 1/32- inch sheets cut to an appropriate size.

3. Composition Determination

The primary method of establishing the compositions of the experimental alloys was by monitoring weight changes

at various stages of the alloy preparation. In addition^ compositions

were checked by wet chemical analysis, electron probe microanalysis,

and lattice parameter measurements whenever these techniques were

appropriate.

4. Temperature Measurement

Temperatures below 1500° C were measured using

platinum versus platinum- 10% rhodium thermocouples. Temperatures

between 1500° C and 2200° C were measured using tungsten versus
tungsten- 26% rhenium thermocouples. These thermocouples were

calibrated against a standard platinum-6% rhodium versus platinum30% rhodium thermocouple up to 1600° C. The standard thermocouple

had been calibrated at the National Bureau of Standards. Secondary

melting point standards were used to check the calibration of the

tungsten versus tungsten- 26% rhenium thermocouples at higher temperatures and to assist in extrapolating a furnace current versus

temperature plot. The secondary melting point standards were high

purity wires of nickel, platinum, rhodium, and iridium. The accuracy

of the temperature measurements was believed to be within ±20° C.

5. Phase Boundary Determinations

a. Solidus - A solidus technique was used

which allows direct observation of a melting specimen and permits

one to obtain its temperature reliably. This technique was suitable

for solidus determinations in alloys which melt congruently or are

close to a eutectic composition. It is also reliable for alloys having

a relatively small temperature increment between their solidus

and liquidus temperatures. For other alloys, the technique of

metallographic examination of quenched alloys to detect incipient

fusion is a more accurate method and was used to supplement the

method of direct observation.

b. Phase Boundaries - The methods used to

locate phase boundaries include metallographic examination, x-ray

diffraction studies, and electron probe microanalysis of equilibrated

two-phase alloys. In each case the specimens were rapidly cooled

or quenched from the equilibration temperature. The methods used

detect changes in microstructure, crystal structure, and chemical

composition which occur when an alloy is equilibrated at various

temperatures or when its composition is varied at a fixed temperature.

4
n. INTRODUCTION

A. Background

This report terminates the investigation of some ten

binary phase diagrams. The research program was a continuation

(Q)

of a previous investigation^ ' of phase relationships among binary

alloys containing a noble metal (i. e. , those metals which are excep -

tionally resistant to chemical corrosion or oxidation, such as gold

and metals of the platinum group).

The existing metallurgical literature concerning

binary or ternary constitution diagrams still contains only a relatively small percentage of alloy combinations involving the noble

metals. The restricted industrial usage of the noble metals due to

their limited availability and high cost may partly account for the

scarcity of constitution diagrams. Another deterrent to such

studies is the prohibitive cost of experimental samples, which

should be of a sufficient number and size to permit measurements

having the desired accuracy. It is probably not surprising, therefore, that the available noble metal constitution diagrams are

sometimes inaccurate and in a few cases may even be considered

to be essentially incorrect.

The importance of obtaining reliable constitution

diagrams is widely recognized among those who are familiar with

their uses. These diagrams may be directly used to predict

equilibrium phase relationships and alloy micros tructures for any

given conditions of temperature, pressure, and alloy composition.

With training and experience it is also possible to use these diagrams

to predict the no n- equilibrium microstructures which are produced

when the alloy is cast, heat-treated, or subjected to various mechanical

operations such as forging, drawing, rolling, etc. The existing

microstructure is a highly important factor in determining the

physical properties of an alloy.

The constitution diagram may be used to reveal other

useful information. It is of considerable importance to know the

temperature at which an alloy begins to melt (solidus temperature).

5
A knowledge of the solidus temperature permits one to estimate

recovery temperatures, recrystallization temperatures, and the

rates of atomic diffusion during annealing treatments. Accurate

constitution diagrams may be used to determine various thermodynamic functions since the slopes of the phase boundaries are

related to free energy relationships between the coexisting phases.

Constitution diagrams are also helpful in studying the kinetics of

phase transformations.

In recent years, there has been a considerable effort

by alloy theorists to develop methods for predicting the extent of

terminal solid solutions and the occurrence of intermediate phases

having various crystal structures. Some attempts have been made

to construct the entire binary constitution diagram on the basis of

theoretical calculations. It appears that significant progress is

being made in this direction, but such studies must continue to

rely on experimental data in order to check the reliability of the

theoretical calculations and to suggest new ways to improve their

accuracy.

B. Experimental Approach

In order to obtain reliable experimental data, it is

necessary to take many precautions during the preparation and

examination of the experimental alloys. Starting materials should

be of high purity and care must be taken to insure that the alloys are

not contaminated during melting or heat-treating. Particular

attention must be given to the problems of temperature measurement

and composition determination since the accuracy of these measurements is a major factor in the production of a reliable constitution

diagram.

Temperatures below 1500° C were measured with

platinum versus platinum- 10% rhodium thermocouples because of the

6
high accuracy and reproducibility associated with this type of

thermocouple. These thermocouples were periodically checked

for accuracy against a standard platinum-6% rhodium versus

platinum- 30% rhodium thermocouple which had been calibrated at

the National Bureau of Standards to 1600° C.

Temperatures above 1600° C are usually measured with

an optical pyrometer, but this method did not yield satisfactory

results in this study due to the formation of metallic films on the

viewing windows. These films formed quite rapidly during the

relatively brief time necessary for a temperature measurement,

particularly when the experimental alloys contained chromium.

It was impossible to make reliable corrections for this film

absorption even by using a two-color pyrometer. Difficulties were

also encountered in obtaining black-body radiation from within the

hot zone of the furnace at a position which would yield a temperature

corresponding to that of the experimental samples. Cylindrical

black-bodies were fabricated of tantalum and designed to produce

the necessary depth to diameter ratios, but it was found that these

objects were unavoidably positioned within rather sharp temperature gradients and were, therefore, not at the uniform temperature

which is essential for such black- body radiators.

In view of these difficulties with optical pyrometry, it

was necessary to consider using high- temperature thermocouples

for measuring the temperature between 1500° C and 2200° C. This

method entails the risk of introducing errors due to thermocouple

contamination by metal vapors. Metal vapors, in some cases, may

even penetrate the refractory ceramics normally used to insulate

the thermocouple, and for this reason it was decided to obtain

thermocouples having an outer metallic sheath. Sheathed thermocouples are commercially available in a variety of sizes and also

7
with several alternative combinations of ceramic insulators and

metallic sheaths. The combination selected as being most suitable

under the conditions of this investigation was a tungsten versus

tungsten-26% rhenium thermocouple insulated with beryllium oxide

and having a tantalum sheath. The size of the thermocouple wires

was 0.005" diameter and the entire assembly was enclosed in a

sheath havii^ a diameter of 0. 040", thus minimizing errors due to

heat conduction along the axis of the thermocouple.

The tungsten versus tungsten-26% rhenium thermocouples

were calibrated against the standard platinum- 6% rhodium versus

platinum-30% rhodium thermocouple to IGOO^'C under the same

experimental conditions used in the actual measurements on the alloy

samples. Simultaneous readings of current in the furnace heating . elements were obtained during the calibration runs so that a plot

of temperature versus current was also obtained. This plot could

be extended to higher temperatures with the help of standard emf

versus temperature data for the tungsten versus tungsten-26%

rhenium thermocouples, together with secondary fixed points obtained by observing the melting of high purity wires of nickel

(1455° C), platinum (1772° C), rhodium (1963° C), and iridium (2447° C).

In addition, many experimental runs were accompanied by an appropriate determination of a secondary fixed point obtained just prior

to the temperature determination. This procedure is explained in

greater detail in reference 6.

Alloy compositions were checked at various stages

throughout this investigation. Considerable reliance was placed

in the use of electron probe microanalysis to determine the compositions of coexisting phases in equilibrated two-phase alloys.

In order to check the reliability of the probe data, however, it is

8
desirable to submit certain alloys to a complete wet chemical

analysis. Alloys selected for this purpose were given a homogenization anneal at a temperature intended to produce a single phase

sample. The compositions selected in each system were usually

close to the equiatomic composition. The alloys were checked for

homogeneity using the electron probe and then submitted to wet

chemical analysis. The result of this chemical analysis was used

not only to check the probe data, but also to check the nominal

composition of the alloy and the feasibility of determining alloy

compositions by monitoring weight losses during melting. Careful

consideration was given to the possibility of alloy contamination

during melting and annealing. Metallographic studies, particularly

at the surfaces of the experimental samples, were quite helpful in

detecting contamination.

Many other difficulties were encountered which posed

a continual challenge to the success of these studies. Some of

these problems are common to all such investigations while others

are unique to a particular alloy system. These problems are discussed in connection with each system.

9
m. VANADIUM-RUTHENIUM CONSTITUTION DIAGRAM

A. Previous Studies

Vanadium- ruthenium alloys have been studied in

(1-3) several previous investigations^ ^ but the information provided

by these studies is still insufficient to permit the construction of

a constitution diagram. A phase having a Cs CI- type ordered

structure has been observed in V-Ru alloys containing 54 to 75

at. % V and a tetragonal distortion of this structure has been

observed very close to the equiatomic composition^ It is not

known whether these structures form directly from the liquid state

or whether they are produced by a solid-state reaction during

cooling. The observation of banded microstructures at the equiatomic

composition suggests a possible martensitic transformation.

Recent studies^^'^^ have revealed sharp changes in

certain physical properties of vanadium- ruthenium alloys in the

vicinity of the equiatomic composition. These anomalies were

observed in measurements of the electronic specific heat coefficient^ , superconducting transition temperature^ , magnetic

susceptibility^^ % NMR Knight shift^^' ^\ and electrical resistivity^^^

It has been reported that the transformation from the cubic

(CsCl-type) structure to the tetragonal structure accompanies these

anomalies at a critical temperature, T^, , which depends sensitively

on the alloy composition^ The following investigation was undertaken in order to help clarify the phase relationships in the vanadiumruthenium alloy system and to provide a constitution diagram.

B. Materials

The vanadium metal used in this study was obtained

from the U.S. Bureau of Mines, BoulderCity Metallurgy Research

10
Laboratory through the courtesy of Mr. T. A. Sullivan. The metal

was in the form of electrolytic dendrites having a nominal purity of

99.95%. A typical analysis of the starting materials is shown in

Table 1.

The ruthenium metal was in the form of powder or

sponge and was obtained from Johnson-Matthey Chemicals Limited.

Its purity was nominally 99. 999%. The impurities detected by

spectrographic analysis are shown in Table 1.

C. Alloy Preparation

Carefully weighed amounts of each starting material

were compressed in a 1/2- inch diameter cylindrical steel mold to

produce small cylinders suitable for arc- melting. Each sample

weighed 10 grams and was arc- melted at least four times in an

atmosphere of 50% helium and 50% argon. Weight losses after

melting were in all cases less than 1% and it was therefore assumed

that the actual compositions of the samples were reasonably close

to the intended compositions. This assumption was later verified

for some of the alloys by electron- beam microanalyses.

D. Temperature Measurement

A platinum versus platinum- 10% rhodium thermocouple

was used to measure temperatures up to 1500° C. This thermocouple

and a standard platinum- 6% rhodium versus platinum- 30% rhodium

thermocouple were used to obtain data on temperature versus

furnace current in a high-vacuum furnace. The standard thermocouple had been calibrated on the International Practical Temperature Scale of 1968 (IPTS 68). Calibration was also made against

11
Table 1

Typical Analyses of the Materials Used in

Preparing the Experimental Alloys

Vanadium Ruthenium

Impurity U.S. Bureau of Mines Johnson- Matthey

C 80

Ca ND 1 Cr 60 ND

Cu 32 <1

Fe : 135 2

H 1 0

Mg ND <1

N . 35

Na ND 2

O 220

Si ND 3

+ Values are given in ppm.

ND Not detected spectrographically.

No other elements were detected spectrographically.

secondary standards such as the melting temperatures of pure

gold, nickel, platinum and rhodium. Temperatures above ISOO^'C

were established by extrapolating the furnace current versus

temperature plot through the points indicated by the secondary

standards and then setting the furnace current to the appropriate

value.

12
E. Thermal Treatments

All alloys were homogenized by annealing at 1600 C for

6 hours except the alloy R^qq^q which was homogenized at 1900°

C

for 6 hours.

The homogenization treatments were carried out

usir^ all the precautions normally taken during an equilibration

treatment and the samples were therefore considered as being

equilibrated at the homogenization temperature. Equilibration

annealing treatments subsequent to the initial homogenization were

carried out consecutively at progressively decreasing temperatures.

The annealing was done in a tantalum- strip resistance

furnace operating at pressures between 5 x 10" and 5 x 10" mm Hg.

The samples were suspended in tantalum buckets using Ta wire.

There was no detectable contamination revealed by subsequent

electron probe studies. Samples annealed in this furnace were

rapidly cooled by turnir^ off the power to the heating elements.

This cooling rate was not always rapid enough to prevent precipitation of lower- temperature phases but the structures formed could

be readily identified. Thus, information on the prior hightemperature phases derived from these structures could be

reliably deduced.

A summary of the equilibration treatments used for the

various samples is shown in Table 2.

13
Table 2

Summary of Equilibration Treatments

for Vanadium- Ruthenium Alloys

Temperatures

rc)

Time Alloys

(atomic %Ru)

1 900 6 hrs 90

1700 3 hrs 80

1600 6 hrs 10, 20,30,40,45, 50, 55,

60, 70,75,78,80,90

1400 1 day 10. 20, 30, 40, 50, 60, 70,

75,78,80,90

1100 1 month 10, 20,30,40,45, 50, 55,

60,70,75,78, 80,90

900 2 months 10, 20, 30,40,45, 50, 55,

60,70,75,78,80, 90

F. Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The solidus temperatures of our alloys were

determined by observations of incipient melting in alloys annealed

at various temperatures above and below the solidus temperature.

Metallo graphic examination was used to detect evidence of melting

for some of the experimental alloys while in other cases it was

possible to locate the solidus temperature with reasonable accuracy

by direct observation of the sample during heating. This latter

method has been described in detail in reference 6 and is a

rather effective technique when working in the vicinity of the

eutec tic compositions or when the solidus and liquidus are

14
separated only by a small temperature interval as is the case for

most of the V-Ru alloys. In most of our alloys, a sharp melting

was observed at a fairly discrete temperature and the solidus

temperature is therefore thought to be located with reasonable

accuracy as indicated on the diagram of Fig. 3.

Losses of metal during a solidus determination

were not excessive. The samples were supported in high-purity

thoria crucibles during the solidus determinations and there \vas

no appreciable reaction between the metal samples and the crucibles

at the temperatures involved.

The temperatures indicated by our thermocouples

during a solidus determination were in good agreement with temperatures deduced by reference to a plot of furnace current versus

temperature provided that care was taken to insure that all

operating conditions corresponded closely to those which prevailed

during our calibration of furnace current versus temperature.

2. Liquidus Temperatures

The liquidus temperatures were not determined

directly but were estimated from the general form of the solidus

curves and the location of invariant reactions. Our metallographic

examination of solidus samples was also helpful in locating the

liquidus through the application of well known principles associated

with the lever law. This was done by estimating the relative amounts

of residual solid and chilled liquid phases in the microstructures of

the partially melted alloys.

3. Other Phase Boundaries

The boundaries of intermediate phases were

determined by metallographic studies of alloys annealed at tempera15
tures \vhich would place them in close proximity to the boundary

being determined. Alloys annealed within two-phase regions were

particularly helpful in determining the location of phase boundaries.

By estimating the relative amount of each phase in the microstructure and then applying the lever rule one can estimate the

approximate location of the phase boundaries. This is particularly

true if one of these boundaries is already known from some

independent prior observations.

In some cases it was possible to determine the

solid-state phase boundaries by quantitative electron probe microanalyses of two-phase (g + a) alloys annealed at various temperatures.

The two-phase alloys, however, were not always reliable for electron

probe microanalyses due to the presence of a fine precipitate in both

the e phase and the a phase regions. These precipitates apparently

form during cooling. For these samples, therefore, the primary

reliance was placed on metallographic studies in locating the phase

boundaries.

G. Experimental Results

1. General

The proposed constitution diagram for the vanadiumruthenium system is shown in Fig. 3. All alloys were examined by

x-ray diffraction or metallographic methods or both and the data

obtained by the different methods were in good agreement. Quantitative

electron probe microanalysis was rather helpful in locating the

boundary of the a and p phase regions but served only to set an upper

limit on the solubilities of these phases at lower temperatures due to

the presence of a fine precipitate in each phase. The electron beam

16
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17
could not be positioned to avoid the precipitate particles due to the

small interparticle spacing. Thus some of these particles are unavoidably positioned within the electron beam and this results in

excessively high contents of one element (V or Ru) being recorded.

Nevertheless, vvith this consideration in mind and utilizing supplementary metallographic and x-ray diffraction data it seems probable

that the phase boundaries have been located with reasonable precision.

2. Ruthenium Terminal Solid Solution - €

The maximum solubility of vanadium in the closepacked hexagonal €- ruthenium terminal solid solution is about 31 at.

% V at 1790° C. There is a significant decrease in solubility at lower

temperatures, however, so that particles of a phase are precipitated

during cooling or after reannealing at lower temperatures (See Fig.

4a and 4b). At 900° C it appears that the maximum solubility of

vanadium in the € phase is slightly less than 20 at. % V. The structure of the precipitated a phase depends on the annealing temperature.

Fig. 5a shows the structure observed in the alloy ^25^^15 after

annealing at 900° C for two months. This structure appears to be

consuming a prior fine dispersoid structure which had apparently

formed when this sample was previously annealed at 1100°C.

Lattice parameter measurements indicate that there

is a slight decrease in aQ, and increase in Cq and an increase in the

c/a ratio with increasing % V. The x-ray data, obtained at room temperature, are given in Table 3. These lattice parameter values are

in reasonably good agreement with those reported by Raub and

Fritzsche^^\

18
208 X (a)

30 at. % V, 70 at. % Ru equilibrated at 1400° C for one day

and rapidly cooled. Large dendrites of GRu containing an

oriented precipitate of ccV phase.

208 X (b)

30 at. % V, 70 at. Ru equilibrated at 1100° C for one

month. Prior fine structure now contains some

cellular growth of 6 Ru and aV phases.

Fig. 4 - Microstructures of vanadium- ruthenium alloys.

19
Table 3

Lattice Parameters of V-Ru Alloys

Alloy Composition Structure Lattice Parameters

a c

iu yu

hep 2 706 4 292

oU

hep 2. 701 4. 295

hep 2.701 4.313

^50^^50 tetr 2.944 3.126

^50^^50 cub (CsCl) 3.003

cub (CsCl) 2.999

cub (CsCl) 2.999

bcc 2.998

^90^^10 bcc 3.017

3. Vanadium Terminal Solid Solution - a

The maximum solubility of ruthenium in bodycentered cubic vanadium is about 55 at. % Ru at 1790°C. The

disordered body-centered cubic structure was observed in all

samples containing up to 20 at. % Ru. The samples beginning

at 30% Ru and extending to the maximum ruthenium contents

showed evidence of atomic ordering. Superlattice lines characteristic of a CsCl-type structure were observed in the x-ray

diffraction patterns of each of these higher Ru-content alloys.

It is not known whether an order-disorder

transformation exists in these samples or whether ordering

20
persists up to the solidus temperature. The smooth continuous

solidus curve containing no sharp discontinuities supports the

view that the disordered body-centered cubic terminal solidsolution is stable at high temperatures up to the maximum

Ru-contents. Raub and Fritzsche^^^ suggest the possible existence

of two-phase regions between the ordered and disordered bcc

phases. However, we have found no evidence for the existence of

such two-phase regions and the lack of convinc ing evidence is the

principal reason why we have not indicated such features on the

proposed diagram.

Tetragonality in the CsCl-type structure was

observed only in our alloy ^qRu^q and in this alloy only when the

sample was annealed at or below 1100°C. According to Chu et al^^^

the cubic to tetragonal transformation in such alloys occurs over a

range of temperatures, Ty^, which vary from about 500° K for

V^,^Rug2 to almost 0°K for V^^Ru^g. Our samples were examined

only at or slightly above room temperature and it is consistent

vvith these results that we observed tetragonality only in our

alloy V^qRu^q. The existence of possible hysteresis effects on

the critical temperature, T^, have been considered^^^ but it was

concluded that any such effects would probably not exceed 0. 5 K.

Obviously further studies will be needed to elucidate the nature of

this transformation.

Quantitative electron-probe microanalysis was

helpful in determining the boundary of the oc- phase region at

1600°C and at 1400° C but the precipitation of a second phase in the

a-phase regions (Fig. 5b and 6a) reduced the reliability of this

method for samples annealed at 1100°C and 900° C. Primary reliance

was placed on metallographic studies for samples annealed at these

lower temperatures.

21
272X (a)

25 at. % V, 75 at. % Ru equilibrated at 900° C for two months.

Prior fine structure is now largely consumed by a courser

structure. Both structures consist of € Ru and aV .

V 'J" '^'r- • '„

160 X (b)

45 at. % V, 55 at. % Ru equilibrated at 1100° C for one month.

Large grains of aV containing an oriented precipitate of

€ Ru phase together with a grain boundary network of 6 Ru.

Fig. 5 - Microstructures of vanadium- ruthenium alloys.

22
272X (a)

45 at. % V, 55 at. % Ru equilibrated at 900° C for two months.

Large grains of aV containing a precipitate of G Ru. Grain

boundary network of G Ru contains a fine precipitate of aV

thus demonstrating a reduced solubility in both phases at this

temperature.

272X (b)

40 at. % V, 60 at. % Ru in the "as-cast" condition.

Eutectic structure of € Ru and aV with fine precipitates

in both phases.

Fig. 6 - Mic restructures of vanadium- ruthenium alloys.

23
4. Eutectic h ^ e + a

The existence of this eutectic reaction was

established by our solidus data and supported by the observation

of a typical eutectic structure in a melted sample having the

composition V^QRug^ (see Fig. 6b). Melting occurs in this alloy

at 1790 ± lO^^C. This is in contradiction to the report by Raub and

(2) Fritzsche^ ' that V-Ru allays in this composition range are molton

at 1700° C. The lower solidus temperatures of Raub and Fritzsche

may perhaps be a result of alloy contamination. We have observed

that the melting points of vanadium alloys can be seriously in

error when crucible reactions are involved. All of our melting

point determinations were done in thoria crucibles in order to

prevent such reactions.

H. Discussion

The proposed V-Ru constitution diagram resembles

the diagram which has been reported for the Nb-Ru system^^^

except that no ordered hep intermediate phase was observed in the

V-Ru system. In both of these systems one observes an ordered

CsCl-type structure and a tetragonal distortion of this structure

which may possibly be the result of a martens itic transformation.

Additional studies are needed in both systems to delineate the

region of stability of these ordered phases and to elucidate their

mode of formation. It is noteworthy that diffusion couples did not

indicate the existence of any two-phase regions separating the

disordered bcc, ordered (CsCl-type) and tetragonal phases for

(1) the Nb-Ru alloys^ '

. High- temperature x-ray diffraction studies,

however, will probably be needed to clarify the mode of formation

of the ordered CsCl-type structure and its tetragonal transformation.

These studies might also be extended to include the Ta-Ru alloys

for which a similar behavior has been noted^^\

24
REFERENCES

1. Greenfield, P. and Beck, P. A., Intermediate Phases in

Binary Systems of Certain Transition Elements, Trans.

AIME 206, pp. 265-276 (1956).

2. Raub, E. and Fritzsche, W. , The Ruthenium-Vanadium

Alloys, Z. Met. 54, pp. 21-23 (1963).

3. Chu, C. W. , Bucher, E., Cooper, A. S. and Malta, J. P.,

Electronic Transitions in Near-Equiatomic VanadiumRuthenium Alloys, Phys. Rev. B. 4, pp. 320-326 (1971).

4. Fliikiger, R. , Heiniger, F. and Miiller, J. , Electronic

Specific Heat and Superconductivity in the VanadiumRuthenium Alloy System, Proc. LT 11, St. Andrews,

Scotland, p. 1017 (1968).

5. Bernasson, M. , Descouts, P. , Donzd, P. and Treyvaud,

A., Magnetic Susceptibilities and NMR Properties of V-Ru

Alloys, _J. Phys . Chem, Solids 30^ p. 2453 (1969).

6. Waterstrat, R. M., The Vanadium- Platinum Constitution

Diagram, Met. Trans. 4^ pp. 455-466 (1973).

7. Hurley, G. F. and Brophy, J. H. , A Constitution Diagram

for the Niobium-Ruthenium System Above 1100° C, _J. LessCommon Metals r, pp. 267-277 (1964).

8. " Constitution of Binary Alloys " (First Supplement) edited

by R. P. Elliott, McGraw-Hill Book Co. , p. 786 (1965).

9. Waterstrat, R. M. and Manuszewski, R. C. , Noble Metal

Constitution Diagrams, NBS Report 10571 (1971).

25
IV. VANADIUM-RHODIUM CONSTITUTION DIAGRAM

A. Previous Studies

The only previous investigation of the vanadiumrhodium alloy system was that of Greenfield and Beck^^^ who

confined their study to four alloy compositions and identified two

intermediate phases. The purpose of the present study \vas to

acquire sufficient information on this system to permit the construction of a constitution diagram.

B. Materials

The vanadium metal used in this study was obtained

from the U. S. Bureau of Mines, Boulder City Metallurgy Research

Laboratory through the courtesy of Mr. T. A. Sullivan. The metal

was in the form of electrolytic dendrites having a nominal purity

of 99. 95%. A typical analysis of the starting materials is shown in

Table 4.

The rhodium metal was in the form of povvder (sponge)

and was obtained from Johnson- Matthey Chemicals Limited. The

impurities detected by spectrographic analysis are shown in Table 4.

C. Alloy Preparation

Carefully weighed amounts of each starting material

were compressed in a 1/2-inch diameter cylindrical steel mold to

produce small cylinders suitable for arc- melting. Each sample

weighed 10 grams and was arc- melted at least four times in an

atmosphere of purified 50% argon and 50% helium. Weight losses

26
Table 4

Typical Analyses of the Materials Used in

Preparing the Experimental Alloys'^

Impurity

Vanadium Rhodium

U. S. Bureau of Mines Johnson- Matthey

C 80 —

Ca ND 1 Cr 60 ND

Cu 32 3

Fe 135 5

H 10

Mg ND <1

N 35

Na ND <1

O 220

Pd ND <1

Si ND 2

t Values given in ppm.

ND Not detected spectrographically.

No other elements were detected spectrographically.

27
after melting were in all cases less than 0. 6% and it was therefore assumed that the actual compositions of the samples were

close to the intended compositions. This assumption was later

verified by electron beam microanalyses of several samples.

D. Temperature Measurement

A platinum versus platinum- 10% rhodium thermocouple was used to measure temperatures up to 1500° C. This

thermocouple and a standard platinum-6% rhodium versus

platinum- 30% rhodium thermocouple were used to obtain data on

temperature versus furnace current in our high- vacuum furnace.

The standard thermocouple had been calibrated on the International

Practical Temperature Scale of 1968 (IPTS 68). Calibration was

also made against secondary standards such as the melting

temperatures of pure gold, nickel, platinum, and rhodium.

Temperatures above 1500° C were obtained by extrapolating the

furnace current versus temperature plot through the points

established by the secondary standards and then setting the furnace

current to the appropriate value.

E. Thermal Treatments

Rhodium-rich alloys containing up to 30 at. % V were

homogenized by annealing at 1650° C for 2 hours. Vanadium-rich

alloys containing up to 10 at. % Rh were homogenized by annealing

at 1700° C for 2 hours. In addition, all of the alloys were annealed

at 1500° C for 2 days before proceeding with any further heat treatments at lower temperatures.

The annealing treatments were accomplished in a

tantalum- strip resistance furnace operating at pressures between

28
5 X 10" and 5 x 10" mm Hg. The samples were suspended in

tantalum buckets using Ta wire. There was no significant Ta

contamination detectable in subsequent electron probe studies.

Samples to be annealed at temperatures up to 1100° C were sealed

in evacuated quartz tubes which had been backfilled with a partial

pressure of argon.

After annealing in the vacuum furnace the samples

were rapidly cooled by turning off the furnace power. Alloys which

had been sealed in quartz tubes were quenched in cold water.

A summary of the equilibration treatments used for

the various samples is shown in Table 5.

Table 5

Schedule of Typical Heat Treatments for Vanadium- Rhodium

Temjperature, °C Time Cooling

1700 2 hrs Rapid cooling*

1650 2 hrs Rapid cooling*

1500 2 days Rapid cooling*

1300 5 days Rapid cooling*

1100 3 weeks Quenched

900 6 weeks Quenched

* Rapid cooling in the vacuum furnace was accomplished

by turning off the furnace power.

29
F. Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The solidus temperatures of our alloys were

determined by observations of incipient melting in alloys annealed

at various temperatures above and below the solidus temperature.

Metallographic examination was used to detect evidence of

melting for some of the experimental alloys while in other cases

it was possible to locate the solidus temperature with reasonable

accuracy by direct observation of the sample during heating.

This latter method has been described in detail^^^ and is a rather

effective technique when working in the vicinity of the eutectic

compositions or when the solidus and liquidus are separated by a

fairly small temperature interval as is the case for most of the

V-Rh alloys. In most of our alloys, a sharp melting was observed

at a fairly discrete temperature and the solidus temperature is

therefore thought to be located with reasonable accuracy as

indicated on the diagram of Fig. 7.

The samples were supported in high- purity

thoria crucibles during the solidus determinations and there was

no detectable reaction between the molten alloys and the crucibles.

2. Liquidus Temperatures

The liquidus temperatures were not determined

directly but were estimated from the general form of the solidus

curves and the location of invariant reactions. Our metallographic

examination of solidus samples was also helpful in locating the

liquidus through the application of well-known principles associated

with the lever law. This was done by estimating the relative

amounts of residual solid and chilled liquid phases in the microstructures of the partially melted alloys.

30
3, Other Phase Boundaries

The boundaries of intermediate phases were

determined by metallographic studies of alloys annealed at

temperatures which would place them in close proximity to the

boundary being determined. Alloys annealed within two-phase

regions were particularly helpful in determining the location of

phase boundaries. By estimating the relative amount of each

phase in the microstructure and then applying the lever rule one

can estimate the approximate location of the phase boundaries.

This is particularly true if one of these boundaries is already

known from some independent prior observations such as electron

microprobe analyses.

In some cases it was possible to determine the

phase boundaries by quantitative electron probe microanalyses

of two- phase alloys (a^ + 3 and 3 + a) annealed at various temperatures. Alloys annealed within other two-phase regions, however,

could not be used for electron probe microanalyses due to the

presence of fine structures or an insufficient amount of one phase.

The fine structures are usually produced by decomposition during

cooling. For these samples, therefore, the primary reliance was

placed on metallographic studies in locating the phase boundaries.

G. Experimental Results

1. General

The proposed constitution diagram for the

vanadium- rhodium system is shown in Fig. 7. All alloys were

examined by x-ray diffraction or metallographic methods or both

and data obtained by the different methods are in good agreement.

31
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32
The use of electron probe microanalysis in locating phase boundaries was limited to alloys containing the 3 phase. All other solidstate phase boundaries were located with sufficient accuracy by

metallographic methods.

2. Rhodium Terminal Solid Solution - y

The extent of the face-centered cubic rhodium

terminal solid solution was established mainly by x-ray diffraction

studies of annealed samples with increasing vanadium contents.

The solubility of vanadium in f cc rhodium extends to about 19

atomic % vanadium and is nearly independent of temperature.

X-ray diffraction patterns of alloys beginning at 20 atomic %

vanadium contain the characteristic "superlattice lines" of the

ordered y' phase. The existence of a two-phase region between the

y and y' phases could not be clearly demonstrated using metallographic or x-ray diffraction methods but its presence is suggested

by a sharp break in the experimental solidus curve at about 20 atomic

% vanadium.

3. Intermediate Phase - y^

The ordered phase (y') possesses a

Cu2Au(Ll2)-type structure. This confirms previous stadies^^'

The degree of ordering is revealed by comparing the relative

intensities of certain lines in the x-ray diffraction pattern of this

structure. Increasing degrees of ordering produce an increase

in the intensity of the so-called "superlattice lines" relative to the

intensities of the "fundamental lines". A visual comparison of

these lines for samples annealed over a range of temperatures from

900° C to 1650° C reveals no pronounced changes in the degree of

ordering. This suggests that the degree of ordering in the y' phase

33
is not strongly temperature dependent. The relative intensities

of the superlattice lines, however, are noticeably reduced at the

compositions V2QRhgQ and V^gRhgg but in these two cases the

intensity reduction is probably due to the co-existence of two phases

at these compositions with accompanying superpositions of some

of the fundamental lines.

X-ray diffraction studies and solidus measurements

indicate that the y' phase (Fig, 8a) extends from about 22 atomic %

vanadium where it forms from the liquid via a peritectic reaction

at about 1740° C to 35 atomic % vanadium at 1615°C where it enters

into the peritectic formation of the € phase. The composition range

of the y' phase is reduced at lower temperatures by slight reductions

in solubility for both rhodium and vanadium (Fig. 8b).

4. Intermediate Phase - €

The intermediate phase e is stable over a narrow

composition range centered at the composition VoRhf-. Its crystal

(A) structure has been determined in a separate investigation^ ' and

may be described as a close-packed array of atoms having a twolayer stacking periodicity and containing both ordered and disordered atomic sites. It apparently forms by a peritectic reaction

at about 1615° C. Its microstructure is characterized by the

presence of extensive microtwinning (Fig. 9a).

5. Intermediate Phase - a«

The intermediate phase is stable over a

composition range from about 40. 5 to 48. 0 atomic % vanadium. It possesses a tetragonal crystal structure in which the close-packed

rhodium and vanadium atoms are ordered in alternate layers along

34
SOX (a)

25 at. % V, 75 at. % Rh equilibrated at 1300 °C for five

days. Twinning in large grains of y' phase.

16 OX (b)

35 at. % V, 65 at. % Rh equilibrated at 900° C for two

months. Large grains of y' phase containir^ oriented

precipitates of C phase.

Fig. 8 - Microstructures of vanadium- rhodium alloys.

35
160X (a)

37. 5 at. % V, 67. 5 at. % Rh equilibrated at 130Cf Cfor five

days. Large grains of € phase containing extensive microtwinning.

40 at. % Y, 60 at. % Rh equilibrated at 900° C for two

months. Large twins and fine deformation markings in

0.2 phase.

Fig. 9 - Microstructures of vanadium- rhodium alloys.

36
SOX (a)

44 at. % V, 56 at. % Rh equilibrated at 900°C for two

months. Twins and fine deformation markings in as phase together with a dispersion of etch pits.

16 OX (b)

47 at. % V, 53 at, % Rh equilibrated at 900° C for two

months. Large grains of phase with fine deformation

lines

,

Fig. 10- Microstructures of vanadium- rhodium alloys.

37
the c-axis and the structure is isomorphous with the CuAu(Ll^)

type. The phase is apparently formed by a peritectic reaction

at about 1560° C. The microstrue ture of this phase contains large

twin bands and other fine deformation markings (Fig. 9b, 10a, 10b).

^ 6. Intermediate Phase -

The intermediate phase is stable between

about 49 and 52 atomic % vanadium and its crystal structure has

been tentatively identified as being isomorphous with the aVIr

structure first reported by Giessen et al^^\ This structure is

orthorhombic and may be considered as a distorted version of the

ordered CsCl-type structure. Our identification of this structure

is based mainly on the position and relative intensities of lowangle lines (Table 6) since at higher angles the lines become

increasingly diffuse. High- temperature vacuum annealing of the

powders did not produce any significant sharpening of the highangle lines, Metallographic studies revealed the existence of

relatively large crystals but each crystal was densely microtwinned (Fig.Ua ). X-ray diffraction data obtained from each of

several individual crystals produced unsatisfactory patterns

suggesting the presence of multiple twinning. We have tentatively

concluded that the phase probably forms by a solid-state reaction

(compare Fig .lib ard 12a). This reaction may occur at about 1400° C

and is probably a peritectoid reaction which we could not suppress

by rapid coolii^. Thus, high- temperature x-ray diffraction

studies will probably be needed to conclusively identify the mode of

formation of this phase. The lattice constants could not accurately

be determined due to the diffuse character of the high-angle lines

but "d-values" obtained from low-angle lines suggest lattice

constants which are close to those reported for aVIr (see Table 6).

38
400X (a)

50 at. ?r V, 50 at. % Rh equilibrated at 1100^ C for three

weeks. Large grains of as phase containing extensive

mic rot winning.

400X (b)

55 at. % V, 45 at. % Rh equilibrated at 1340° C and rapidly

cooled. Fine structure of a.3 and 3 phases formed from a prior matrix of ai and a trace of 8-

Fig. 11 - Microstructures of vanadium- rhodium alloys.

39
SOX (a)

55 at. % V, 45 at. % Rh equilibrated at 1100° C for three

weeks. Matrix is aa phase and contains oriented,

elongated islands of 3 phase.

680X (b)

57. 5 at. % V, 42. 5 at. % Rh in the "as-cast" condition.

Fine eutectic structure composed of g and decomposed

ai phases.

Fig. 12 - Microstructures of vanadium- rhodium alloys.

40
7. Intermediate Phase -

The existence of a high- temperature phase on

the vanadium- rich side of the equiatomic composition is suggested

by metallographic studies in which we observed the finelydecomposed structures (FigJlb,13a ) which are characteristic of a

eutec toid decomposition. Metallographic examination of alloys

quenched from various temperatures indicated that the eutectoid

temperature was located at about 1325° C (see Fig. lib and 12a). We

were unable to retain the high- temperature phase in alloys which

had been rapidly cooled from above this proposed eutectoid temperature and we were therefore unable to identify the crystal

structure of the phase.

8. Intermediate Phase -

p

Previous studies have established the

existence of a phase having the A15 (p-W) structure in the V-Rh

system at the "ideal" stoichiometric composition (A^B). We have

confirmed these observations and, in addition, we find that this

phase is stable over a substantial composition range extending

toward the rhodium-rich side of the "ideal" composition. The p phase apparently forms by a peritectic reaction at 77 atomic %

vanadium and at 1730° C. On the rhodium- rich side it enters into

a eutectic reaction with a^^ at 1520° C. There is a slight reduction

in solubility for both vanadium and rhodium at lower temperatures

where the stability range extends from about 63. 0 to 74. 5 atomic %

vanadium. A similar range of composition has been reported for

this structure in the V-Ir alloy system where the superconducting

transition temperatures have been reported to increase significantly with increasing iridium contents^'^^ A similar effect has

now been observed in our V-Rh p-phase samples. The p-phase

alloy VpcRh^c was found to have a superconducting transition

41 I
400X (a)

60 at. % V, 40 at. % Rh' equilibrated at 1500° C for two

days. Large islands of b phase interspersed with a

fine structure of decomposed phase.

400X (b)

80 at. % V, 20 at. % Rh equilibrated at 1500° C for two

days. Fine two- phase structure consisting of 3 and aV.

Fig. 13 - Micros true tu res of vanadium- rhodium alloys.

42
temperature of 1.075K^ ' whereas the alloy Vr^gRhgg remained

in the normal state at 0. 015K^^^. The extent of the p-phase region

in the V-Rh system was established primarily by quantitative

electron-beam microanalyses of alloys annealed and quenched from

various temperatures.

9. Vanadium Terminal Solid Solution - a

The body- centered cubic vanadium solid

solution was found to dissolve up to 18 atomic % rhodium at 1730° C

where it enters into the peritectic formation of the p phase. The

solubility of rhodium in the a solid solution decreases to about

14 atomic % rhodium at lower temperatures. Two-phase (a + p)

alloys reveal an unusually fine, uniform structure (Fig.13 b).

10. Invariant Reactions

All invariant reactions found or assumed in the

V-Rh system are listed in Table 7. Many of these were established by solidus measurements. Solid state reaction temperatures

were determined by bracketing the reaction temperature with

annealing runs and then examining the x-ray patterns and microstructures of these alloys. In some cases the reactions were only

postulated from circumstantial evidence since it was not possible

to detect them by the methods used in this study.

In many cases it was quite helpful to first

establish the solid-state phase boundaries by x-ray microprobe

analyses and by examining the microstrue tures of two- phase

alloys. These solid-state boundaries were then extrapolated to

the reaction horizontals established by the solidus measurements.

Some applications of the lever law were necessary to establish the

43
eutectic or peritectic compositions. This was accomplished

by metallographic examination of the samples used for the

solidus measurements. The accuracy of such methods is

difficult to estimate but errors vvould probably not exceed ±2 at. %.

Errors in temperature measurement are probably ±10° C.

H. Discussion

The vanadium- rhodium constitution diagram which

emerges from the present investigation is similar in appearance to the vanadium- iridium constitution diagram tentatively

presented by Giessen et aP . At least four of the six intermediate phases observed in the V-Rh system (y', a«, and p) are

structurally equivalent to phases observed in the V-Ir system^ ' [a-VIr^, 3-(Vj ^Ir^)Ir, a-VIr, and a-V^Ir, respectively]. These

four phases occur over similar ranges of composition in both

systems. Other similarities may also exist but the data so far

available on the V-Ir system are inadequate for a complete

description of the constitution diagram and more data are needed,

particularly at lower temperatures.

One may also observe similarities between the V-Rh

constitution diagram and the Nb-Rh constitution diagram presented

by Ritter et al^^^\ In both systems one observes a high temperature phase (a^) of unknown crystal structure \Vhich decomposes

eutectoidally on cooling. In addition three other V-Rh intermediate

phases (y', and p) have structural counterparts in the Nb-Rh

system [aNbRhg, a2(NbQ ggR^Q Q^)Rh and aNb^Rh, respectively].

The € phase (V^Rhg) apparently resembles the phase a^fNbgRhg)

since both structures contain close-packed atom layers having a

mutually equivalent type of atomic ordering^^' The t\vo

44
structures apparently differ only in the stacking of these layers^

which follows a two-layer sequence in V^Rhg compared to a ninelayer sequence in Nb^Rhg. Another interesting feature of these

two systems involves the existence of a broad sigma phase region

in the Nb-Rh system which is apparently replaced by a similarly

broad A15-type phase region in the V-Rh system. A plot of the

atomic volumes of the V-Rh phases is sho\vn in Fig. 38. It involves a slight negative departure from a linear relationship

similar to the behavior observed in both V-Ir^^^ and Nb-Rh^^^^

systems. Crystallographic data and lattice parameters for phases

forming in the V-Rh system are summarized in Tables Sand 9.

45
Table 6

Relative Intensities (I) and Interplanar Spacings (d) for Phase

Lines Compared with Equivalent Values Reported for

the Phase aVIr^^^

g^VRh aVIr

hkl d*

110 - - - - 020 3. 325 w 3.384 20

200 2.888 w 2.887 20

001 2.771 w 2.805 <20

111 2.340 w 2.361 10

220 2.183 s 2. 200 100

021 2.136 s 2.154 60

130 2.073 w 2.102 40

201 2.005 s 2.017 80

310 1. 853 vw 1.857 10

221 1.718 mw 1.726 50

040 U666 m 1.693 30

131 1.682 30

31 1 330 1.467 <10

240 1.463 <10

400 / 041 i 1.449 mw 1.449 20

002 1*392 m 1.400 20

112 ) 420 i 1.329 w 1. 330 20

150 1. 317 10

(Table 6 continued)

46
Table 6 (continued)

hkl

VRh aVIr

li d* I d I

331

241 1. 284 m 1.294 60

022

401 1.253 w 1.287 10

202 _ 1.260 < 5

421 m

1.202 30

151 1.192 20

222 1.176 ms 1.181 35

132 1. 157 vw 1. 166 10

510

060

312 1.113 vw 1.118 <5

350 1.098 w 1 . 108 15

440 1.070 w 1.099 15

* Lines of a^VRh become increasingly broad with

increasing Bragg angle and the accuracy of the values

reported probably does not exceed ±0.02 L On the

basis of the d values above, we would estimate that

the lattice constants for a^VRh are:

a = 5. 78 b = 6.65 c = 2. 78

This may be compared with the lattice constants reported

for aVIr vvhich are:

a = 5.791 b = 6.756 c = 2.796

(5)

47
B

+->>> COB

.2T3OK

.2>

(DO

-4-> O

a» K

-t->

.1-1 ua

COuaBQ QJ O 0) ^ CO o

-t-> 0)

0) o o.1-1 +-> o

0)ao

T3

0) CU 0)

-i-> -i-> cti i-H i-H I—

3 3 5

CO CO CO o o o

Ph fl| S 5So

CO

.r-l 1—1o

COo o CO CM

COd

CD B

-•-> OOOo

CD

;3O

CD CO

CO

CDUBB i ^CD

J

*-l

a;

H

0)a

CO

UU

CO CO

CD CD

aaa

u

CO

CO

CD

.1-1 73

COo

CO

0)

.1-1

CO

O .^5O

X2a

CO 73

CO CO

.1-1 .1-1 1—1 r—

O O

CO CO

CD^ be P '-^ 'O oi

.1-1 -(-)

CO ^ o

in o o CO

lO

CD

in ^ 2 CD I> S3

I—t r—I r—I 1—

a ^ beor—

CO

en ^+

05CMO• C30

1—1 >

05

COoa

II

CM*

LO• o • ITSCO d

II CO d+

"o

00 CM* l> in i>

CO d> CM

ifi

d cn. >- + II 1—

CM d

II

d "o

IICM IIII OS CD*

d+ + 4- + 1—

d

d oO .-5

CD

CD 13 ^ CM o o o o

.1-1

•1-1 .1-1 •1-1 .1-1 s TS

4-> -l-> -l-> •1-1 o o CJ o O

0) CU cu o

;h ;h ;h ;h

u

(1)

0) (D CD CD ;=( Ph Ph Ph p^ ft W 1—

1—1—

1—1—1—

CO LO CD D- 00 1 1 1 1 1 1 1 1 1 1 1 1

48
r:

O CO S B O -r^

•r-l ^ I—

O

z;

0)a.

CD BapoO

cuoa

0)

-l->

CO

i-Hd*-> muO

0)

-i->

•r-l s o O in

05*0

CO o mo ^ o ^ o ^ lO

Oc^ ^ in s§CO o in '-• in CM* CM in

00oCO uO

inc

COa

CD « o

•i-i

CO > B

S3Oc

X3>

rH-* oU

-MEh

X2

CM

I—

CD

x:

00

COd>

CM d

-t-> o

CO

00>I dOI

rH

WQ

.aoo

x:

-t-> uO

x:>

CO d

CD o

1—1 in

CD

X3oa

a

CM

O s a

oo

So"

CO

in ^

CD* CO

CO O<

CM

COaI ao

x:uo

x:O

X3

inCMd

in ;C . CO >°>O o.1-1 O

CO

•C>

a,§u

CDa

03a,o

CO

-(-> %

(U

CDO.1-1 o

x:o

0)

CDU

CD§o

(DU

CD

i-H

oi

1—1

•r-l d

ci

0) X3H

49
§

o•i-l -M: o

O

>:t< lO Tj< CO

Oi Oi Oi

uH

-sd

in

cuCO

ft

cu

•i-i 73

0)a

B

.3O

x:

I a

B

Pi >

o

Cfl

-t-> CU

ft

<uo

05 t- OS O

00 05 in OS CSl C- lO CD ITS CO

CM* CO* CO* CO* -"J^^*

05O 00

tOS 00

tr- 00

tACO

• CO

in

00 OS o

00 00

CO

ITS O

\n CM 00

00 00 OS CO00 OS OS oo

oo

CO ^ ^ Lfscococoincocococococo

COu

I

u

ei ft

Q)O

•J

O)

Cfl OS

ft oa ca cn

co csi CM d d c5 d >- > >^ :^ >^ >-

ooooooooinoooooo

irs in o ,—I CM COirs CO O CO CD o

in in in CD

CMCD o in CM in o in I> 00 00 OS <7S

50
REFERENCES

1. Greenfield, P., and Beck, P. A., Intermediate Phases in Binary

Systems of Certain Transition Elements, Trans. AIME

206 , pp. 265-276 (1956).

2. Waterstrat, R. M., The Vanadium- Platinum Constitution

Diagram, Met. Trans. 4, pp. 455-466 (1973).

3. Dwight, A. E. and Beck, P. A., Close-Packed Ordered

Structures in Binary AB^ Alloys of Transition Elements

Trans. AIME, 215, pp. 976-979 (1959).

4. Waterstrat, R. M., and Dickens, Brian, The Crystal

Structure of V^Rhg, J. Less-Common Metals 31,

pp. 61-67 (1973).

5. Giessen, B. C. , Dangel, P. N., and Grant, N. J.,

New Phases in the Vanadium- Iridium System and a Tentative Constitution Diagram, _J. Less- Common Metals

13, pp. 62-70 (1967).

6. Zegler, S. T., and Downey, J. W. , Ternary CrgO- Type

Phases with Vanadium, Trans. AIME 227, pp. 1407-1411

(1963).

7. Cox, J. E., Bostock, J., and Waterstrat, R. M. , Superconducting Properties of A15 Phase V-Ir Alloys.

Proc. LT 13, Boulder, Colorado (1972).

8. Cox, J. E., and Waterstrat, R. M. , unpublished results

(1973).

9. Blaugher, R. D. , Hein, R. E., Cox, J. E., and Waterstrat,

R. M. , Atomic Ordering and Superconductivity in A15

Compounds, J. Low- Temp. Physics 1, pp. 539-561 (1969).

' 51
Ritter, D. L., Giessen, B. C, and Grant, N. J.,

The Niobium (Columbium)- Rhodium Binary System, Part I: The Constitution Diagram, Trans. TMS-AIME 230,

pp. 1250-1259 (1964).

Ritter, D. L. , Giessen, B. C, and Grant, N. J.,

The Niobium (Columbium)- Rhodium Binary System, Part II: Crystal Structure Relationships, Trans. TMS-AIME 230,

pp. 1259-1267 (1964).

52
V. NIOBIUM-OSMIUM CONSTITUTION DIAGRAM

A. Previous Studies

Niobium- osmium alloys have been the subject of

several previous investigations^ ' but there has been no attempt

to construct a constitution diagram for this system. In fact, one

observes a disagreement among the various investigations

regarding the location of some phase boundaries. Nevertheless,

the existence of at least three intermediate phases in this system

seems to be well-established. The present study represents an

attempt to obtain sufficient data to permit the construction of a

constitution diagram.

B. Materials

The niobium metal used in this study was obtained

from Leico Industries, Inc., in the form of powder (-325 mesh)

having a nominal purity of 99. 8%. A portion of this material was

melted in our arc- furnace and then submitted to spectrographic and

vacuum-fusion analyses with the results shown in Table 10.

The osmium metal was also obtained as powder

(-325 mesh) from Engelhard Industries, Inc. and had a nominal

purity of 99. 9%.

A portion of each of these metals was arc- melted and

then submitted to spectrographic analysis with the results shown

in Table 10.

53
Table 10

Typical Analyses of the Materials Used in

Preparing the Experimental Alloys '

Impurity

Niobium Osmill m

Leico Industries Inc Enffelhard Tndnstrips Tnp

Al < 300 < 1 0

Au < 1 0

c 53

Ca < 10J. w <^ 1X 00\J\J Cr * <30

Cu * <10

Fe < 1000 <30

H 1 5

Ir < 300

Msr < 100 ^ t-l V/ N 95

Ni ND <10

\j OODD

Pt <100

Rh <100

Si <1000 <30

Ta * V ND <10

W ND <30

* Not determined due to niobium interferences.

+ Values given in ppm.

ND Not detected spectrographically.

No other elements were detected spectrographically.

54
C. Alloy Preparation

In preparing the experimental alloys we obtained

helpful guidance from previous studies of the Mo-Os system^^^

(5) and the Nb-Ir system^ \ For example, special precautions

were taken to insure that no residual unmelted portion of the

osmium metal remained after arc- melting. This was accomplished by a thorough mixing of accurately weighed portions of

niobium and osmium powders. The mixture was compressed in

a 1/2- inch diameter cylindrical die of hardened steel to produce

small cylinders suitable for subsequent melting. The cylindrical

powder compacts were then sintered and degassed in a high5 6

vacuum furnace operating at 1600° C and between 10" and 10" mm Hg.

Previous experience in sintering pure Os had suggested that a

weight loss of about 0. 5% can be expected during this procedure and

we, therefore, added a small excess of Os to each alloy compact in

order to compensate for this small loss and thus produce an alloy

which closely approximates the intended nominal composition.

When the sintering had been completed, the alloys were removed

from the high-vacuum furnace and each alloy composition was then

melted at least four times in an arc- furnace under an inert atmosphere of high-purity 50% argon and 50% helium. The arc current

was maintained at a high level during melting (-600 amperes) to

insure complete melting of each alloy. Subsequent visual and metallographic studies indicated that all samples had been adequately

melted and electron microprobe analyses of several alloys revealed

that no significant contamination had occurred. Each alloy sample

weighed about 20 grams and was given a homogenization anneal at

high temperatures before sectioning the sample for internal examination and subsequent lower -temperature annealing.

55
D. Temperature Measurement

Temperatures up to 1500° C were measured with a

platinum versus platinum- 10% rhodium thermocouple. This thermocouple was calibrated with a standard thermocouple of platinum6% rhodium versus platinum-30% rhodium which was itself

calibrated at the National Bureau of Standards on the International

Practical Temperature Scale of 1968 (IPTS 68). Both of these

thermocouples were used to obtain data on temperature versus

furnace current in our high- temperature vacuum furnace. The

furnace was in this case equipped with a tungsten- mesh heating

element and tungsten heat shields rather than the usual tantalum

fixtures so that temperatures in excess of 2500° C could be obtained.

Temperatures from 1500° C to 2200° C were measured with a tungsten

versus tungsten-6% rhenium thermocouple. At all temperatures, and

especially at temperatures above 2000° C, it was necessary to

supplement our temperature versus furnace current plot with data

obtained from the observed melting points (IPTS 68) of certain

secondary standards. The secondary standards used in our calibration procedures were high- purity wires of gold (1064° C),

nickel (1455° C), platinum (1772°C), rhodium (1963° C), iridium

(2447°C), niobium (2472°C) and molybdenum (2625°C). In most

cases the standard wires were placed in thoria crucibles using a

technique described previous ly^^^. At about 2500° C, however, it

appeared that there may have been some reaction between the thoria

crucibles and their tantalum holder and, therefore, the niobium

and molybdenum wires were simply suspended from fine tungsten

wires extending part way into the hot zone of the furnace. Each

standard wire was formed into a small (~l/8-inch diameter) loop

at its tip to facilitate the observation of melting. The wires always

melted first at this tip since the tip was accurately placed in the hot

zone at the position of maximum temperature.

56
The accuracy of these procedures, however, can be

no better than the accuracy to which one can specify the melting

points of the secondary standards. At present it appears that an

accuracy of ±10° C would be the best that could be claimed for our

Nb and Mo wires which had a nominal purity of at least 99. 9%.

Nevertheless, future work may produce more reliable values for

wires of this purity and if more reliable values are obtained it should be possible to make appropriate corrections to our data.

There is certainly an urgent need for accurate secondary temperature standards above 2000° C since many experiments are not

conducive to the use of ideal "black- body" radiation cavities either

because of interference with other objects in the hot zone or

because the hot zone does not contain a region of reasonably

constant temperature large enough to accommodate the "blackbody" cavity. In such cases the use of radiation pyrometry

techniques will suffer a substantial loss of accuracy even though

the accuracy and sensitivity of some present photo-electric devices

are quite good. The tungsten- rhenium thermocouples are, at best,

only a partial solution to this problem since substantial changes in

their calibration can occur during a measurement either as a

result of contamination or because of strains introduced by thermal

expansion and contraction during temperature cycling. Even with

these thermocouples, however, secondary standards would be very

helpful in improving measurements above 2000° C.

E. Thermal Treatments

Alloys containing 80% Os or more were homogenized

by annealing at 2420° C for 6 hours. All of the alloys were annealed

at 2000° C for 12 hours prior to any annealing at lower temperatures

All annealing was performed in a tungsten- mesh

- 5 -6

resistance furnace operating at pressures between 10 and 10 mm

57
The samples were suspended in Ta buckets lined with Nb sheet

and suspended with Ta wire. Samples to be annealed at 1100°

C

were sealed in evacuated quartz tubes which were back-filled with

a partial pressure of argon gas.

' Samples annealed in the vacuum furnace were

"radiation-quenched'' by turning off the furnace power. Alloys

sealed in quartz tubes were quenched in cold water. The cooling

rates appeared to be adequate for retaining structures formed at

the respective annealing temperatures except for the a phase

containing over 68 at. % Nb which always transformed to the b phase

during cooling. A summary of typical annealing treatments used

in this study is given in Table U.

Table U

Summary of Typical Ecmilibration Treatments

for Niobium- Osmium Alloys

Temperature Time Alloys

°C (atomic %Nb)

2420 6 hrs 10, 20

2000 12 hrs 10,20, 30,36,45,50,56,65,

70,75, 80, 90

1600 1 week 10,20,30,36, 45, 50,56,65,

70, 75,80,90

1100 6 weeks 10, 20,30,36, 45,50,56,65,

70, 75, 80, 90

58
F. Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The solidus temperatures of the niobium- osmium

alloys were determined mainly by direct observation of small

(~1 g) samples during heating. The samples were tied in place at

the end of a piece of fine (. 005") tungsten wire or, particularly at

the highest temperatures, with thicker (.010") tungsten wire. The

tungsten wires were then suspended so that the sample would be

located at the point of maximum temperature in the hot zone.

Standard wires of Pt, Rh, and Nb were also suspended in this

manner in preliminary experiments which established that the

methods was capable of yielding accurate melting point data.

The method was surprisingly successful. In

many cases the sample melted to form a small molten droplet

which remained fixed in the loop of tungsten wire presumably with

some support provided by surface tension of the liquid droplet as

it wetted the tungsten wire. One might expect that at these temperatures (>2000°C) there would be significant diffusion and

chemical reaction between the niobium- osmium alloys and the

tungsten \vire but subsequent examination of the wires and the

samples revealed little evidence of any reaction although the wire

was indeed "wetted" by the molten alloys. The wire in contact

with the sample essentially retained its original size and shape,

as far as could be determined by visual observation, although

recrystallization of the wires at these high temperatures did produce embrittlement. The absence of any significant reaction

between the wires and the samples suggests that the method has

probably yielded reasonably accurate solidus temperatures for the

59
niobium- osmium alloys. One must remember that the temperature is recorded at the first sign of melting and that reactions

in the solid state are likely to be even more restricted than

reactions with the liquid. Furthermore, it appears that, in this

system, melting occurs rather sharply in most cases due to the

small interval between the liquidus and solidus temperatures.

— 2. Liquidus Temperatures

The liquidus temperatures were estimated from

the general form of the solidus boundaries and from the location

of the invariant reactions. In some cases metallographic observations on samples used for solidus determinations were helpful in

suggesting the location of the liquidus.

3. Other Phase Boundaries

The boundaries of the various phases were

determined by metallographic and x-ray diffraction studies together with quantitative electron-probe microanalyses of the

phases coexisting in two-phase alloys. The relatively narrow

two- phase regions occurring in this system permit the location of

phase boundaries with unusual accuracy through use of the wellknown "lever rule". One first estimates the relative amounts of

each phase coexisting in a two-phase alloy and then, by

independent knowledge of one of these phase boundaries, it is

possible to estimate the approximate position of the other boundary.

A fortunate choice of alloy compositions lying very close to the

phase boundaries in this system greatly facilitates this process

since one phase boundary is thus already fairly well established.

The accuracy of the phase boundaries determined in this study is

estimated at about + 1 atomic %.

60
61
G. Experimental Results

1. General

The proposed constitution diagram for the

niobium- osmium system is shown in Fig. 14. The number of

alloy compositions represented by our experimental samples is

relatively small but these were, fortunately, located in such a

manner that considerable information on the location of the phase

boundaries was obtained. The proposed diagram is, therefore,

considered to be accurate to about ± 1 atomic % and ± 20° C. The

crystal structures for each intermediate phase are given in

Table 12 and lattice parameters are given in Table 13.

2. Osmium Terminal Solid Solution - e

The close-packed hexagonal osmium terminal

solid solution dissolves up to about 27 at. % Nb at 2270°C. An

increasing Nb content in this phase is accompanied by an increase

in the lattice parameters (a^ and c^) with a nearly constant axial

ratio (c/a). The solubility limit is only moderately temperature

dependent. Alloys containing more than 27 at. % Nb (Fig 15a) show

that the € solid solution coexists with an intermetallic compound

(v phase).

3. Intermediate Phase -y

The X phase was found to possess a cubic

fl 3 a-Mn (A12)-type structure, in confirmation of previous studies^ '

It extends over a significant composition range from about 35 to 51

at. % Nb and apparently forms via a perifcectic reaction at

about 2270° C. Although previous studies suggested that atomic

ordering was present in this structure, there has been little

62
fe5

O CO

O -r-l

CO

o

5

o

-l-> COa

g

O

o

a

<D

-»-> CQ

(/2

CO

(D

• r-l T3

o o o -M O

CO ^

00

CQ

T3

T3

U

O

a

CO

1

o•i-i 5

CD ^

o O OO

IT)

O

CO

cnhCOQ

o

'oD

H

CD 0

—1 <^

00

10

T3

CO 00

CO

ft

o

o•r-l u

£0h

63
o

1—

ID

00

Oi

in

00

LO

a

CO o

I a

.2o

o

COo

COu

-l-> s

O

CO O

CO ID o o

in

in 00 —» CO

CO

'if

o

CO

o

cu

CO

ft

CO 1—

(

CO lO 00 tr- O CD 1—

(

CM CM CD lO CO CO 05 CM t- CO CD CD in CM CM CM 1—1 1—

00 00 t- t- CD t- tCO CO CO in uo C35 Oi Oi CJ5 05 CM CM

CO CO oa D b X X o o

d d

CO o

in

CM CM ineo O

in

in

in CD o

00

o

05

64
Table 14

Observed and Calculated Relative Intensities

for NbggOSg^ X- Phase

hkl Iobs Disordered Orderedt

110 - 0 1 200 - 1 0

211 m 5 110

220 w 7 44

310 vw 2 10

222 vw 8 14

321 - 7 4

400 m 48 60

330 i

) 41l'1 vs 1 AAA i 000 1 000

420 vw 5 10

332 s 278 441

422 m 101 107

431)

510 i ms 140 262

521 12 24

440 w 2 6

433 1 530 '[ mw 11 15

442 1 600 '

!

" 14 11

532 j 611 '

1 mw 20 13

620 w-^ 3 7

541 4 5

(Table 14 Continued)

65
Table 14 (Continued)

hkl I , Disordered Orderedt

622 m 27 32

631 vw 7 4

444 ms 71 85

543

550 s 83 112

710

640 — 0 0

552

633 vs 308 326

721

642 s 54 109

730 ms 36 51

732

651

s 54 72

800 0 0

811

741 m 15 18

554

820

644 1 m 24 21

653 m 17 22

822

660

* vs 98 132

831

750 m 31 26

743

/ (Table 14 Continued)

66
Table 14 (Continued)

Icalc

hkl I^^g Disordered Ordered"

662

752

840

910)

833 t 842

921

761

655

664

930

851

754

w

mw

ms

w

ms

w

ms

9

11

1

31

9

25

19

59

7

14

1

64

14

63

18

69

+ The model chosen was based on the composition

NbggOSg^ and therefore one must distribute

approximately 21 Nb atoms among the 58 atomic sites

in this structure. The large Nb atoms were

assumed to be distributed so that 10 Nb atoms

completely fill the sites of CN 16 [positions 2(a)

and 8(c)]. The remaining 11 Nb atoms were

assumed to partially fill the sites of CN 14

[position 24(g)] with Os atoms occupying the re- mainder of these sites. The sites of CN 12

[position 24(g)] are assumed to be completely

filled with Os atoms.

67
160X (a)

30 at. % Nb, 70 at. % Os. Annealed at 2000°

C

for twelve hours. Globules of y phase in a matrix

of e Os.

SOX (b)

50 at. % Nb, 50 at. % Os. Annealed at 1600°C

for one week. Grain boundary network of

Sigma phase surrounding large grains of y phase.

Fig. 15 - Microstructures of niobium-osmium alloys.

68
inl'ormation available on the nature and extent of this atomic

ordering. We have, therefore, calculated the relative intensities

to be expected for an x-ray diffraction pattern from what we

considered to be the most likely ordered structure. This involves

placing the larger Nb atoms preferentially in sites having the

highest coordination number (CN 16). Observed and calculated

intensities are presented in Table 14. The relative intensities

have been calculated for both ordered and disordered models. The

disordered model is clearly ruled out by significant differences

between observed and calculated intensities. The ordered model

vvhich we chose, however, eliminates these differences and

produces a reasonably good agreement between calculated and

observed intensities. Thus it appears that the proposed model for

the atomic ordering may be substantially correct. The x phase

enters into a eutectic reaction with another intermetallic compound

(cT phase) and a small tvvo-phase region exists near the equiatomic

composition (Fig. 15b ).

4. Intermediate Phase - ?

The existence of a Nb-Os sigma phase was first

reported by Knapton^"'^^ The sigma phase region extends from

about 55 to 73 at. ? Nb at 1975°C and it apparently melts

congruently at about 2200' C. It possesses a tetragonal (D8, )-type

(I) structure^ ^ similar to that of beta-uranium. The atomic ordering

(1) in this phase has been determined in a previous study^ \ The Osrich boundary of the i phase and the Nb-rich boundary of the

X phase were located by microprobe analysis of the sample

Nb^QOs^Q. This sample was essentially single phase (x) except

for a surface layer which had apparently been depleted in Os during

annealing and which, consequently, had transformed to the sigma

phase. Microprobe analyses were made close to the boundary

69
between the y phase interior and the a phase surface layer and on

both sides of this boundary.

5. Intermediate Phase -

g

^ The beta phase region extends from about 70. 5

to 75 at . % Nb and this phase forms via a peritectoid reaction

at 1975°C. It possesses the cubic A15-type structure in which

partial atomic long-range ordering occurs. The extent of the

atomic ordering has been reported in a previous study^^^. This

(Q) phase was apparently first discovered by Geller et ar It coexists with the a phase (Fig. 16a) and the aNb phase (Fig. 16b).

6. Niobium Terminal Solid Solution - a

The body-centered cubic niobium solid solution

may contain up to 19 at. %Os at 2175°C but a substantial decrease in solubility to about 14 at. % Os occurs at lower

temperatures. This was established mainly by x-ray lattice

parameters measurments since the presence of fine precipitates

in this phase produced erroneous x-ray microanalyses. The

addition of osmium to the cc phase produces decreasing liquidus

and solidus temperatures culminating in a eutectic reaction with

the Sigma phase at about 2175° C.

7. Invariant Reactions

There are five invariant reactions in the niobiumosmium binary system as follows:

A eutectic reaction occurs at 217 5° C;

h = a + a (See Fig.l7 a)

70
SOX (a)

70 at. % Nb, 30 at. % Os. Annealed at 1600°

C

for one week. Matrix of p phase containing

globules of Sigma phase in which some precipitation of secondary p has occurred during cooling.

BOX (b)

80 at. % Nb, 20 at. % Os. Annealed at 1600°C

for one week. Globular clusters of a Nb in a matrix of p phase.

Fig. 16- Microstructures of niobium-osmium alloys.

71
160X (a)

75 at. % Nb, 25 at. % Os, Melted primary sigma

phase with interdendritic eutectic of sigma plus

a Nb phases.

160X (b)

75 at. % Nb, 25 at. % Os. Annealed at 2000°C

for twelve hours, a Nb globules in a matrix of

sigma phase.

Fig. 17 - Microstructures of niobium- osmium alloy

72
A peritectoic reaction occurs at 1975 C;

a -h a = 3

Congruent melting of the sigma phase occurs at 2200° C;

L = a A eutectic reaction occurs at 2120° C;

L = X + a A peritectic raction occurs at 2270° C;

e + L = X

The existence of most of these invariant reactions

was established by the solidus measurements and by metallographic examination of the solidus specimens themselves. The

peritectoid reaction temperature was established by annealing the

alloy Nb,^gOs25 at various temperatures and noting any microstructural changes (See Fig. 16b and 17b).

H. Discussion

The niobium-osmium constitution diagram (Fig.l4) is

dominated by the presence of the three intermediate phases y, a and p. These phases share some common crystallographic

features. Each structure seems to contain atoms of differing

size arranged in a dense packing such that only tetrahedral interstices are formed and giving rise to some characteristic atomic

polyhedra (Kasper polyhedra) with relatively high coordination

numbers^^^^. The structure of the X phase contains Kasper polyhedra having apparent coordination numbers (CN) of 12, 14 and 16.

The structure of the sigma phase contains Kasper polyhedra with

the coordination numbers 12, 14 and 15 and the structure of the

p phase contains Kasper polyhedra based on the coordination

numbers 12 and 14. These polyhedra are sometimes distorted,

apparently in response to the packing requirements, so that the

73
actual coordination numbers, in some cases, may not be clearly

defined. Nevertheless, these structures contain only tetrahedral

interstices in contrast to the common face-centered cubic or

hexagonal close- packed structures in which there are both tetrahedral and octahedral interstices Observed variations of the

interatomic distances in the a and 3 phase structures suggest

that a primary factor in their formation is the availability of

atomic sites of high coordination^^^^

Atomic ordering is commonly observed in these structures and, as a rule, one finds elements with the largest atomic

size in the sites of highest coordination number (15 or 16) whereas

the smaller atoms seem to prefer atomic sites with the lowest

coordination number (12). Atomic sites with intermediate

coordination numbers (13 or 14) contain varying mixtures of both

large and small atoms and are often associated with the significant

composition variations occurring in these phases.

The extent of atomic ordering in the various

crystallographic sites has been determined for each of the

niobium-osmium intermediate phases and this data is summarized

in Table 15. Obviously the extent of filling in each site is

determined partly by the over-all composition of the phase but there

is also a clear indication that the larger niobium atoms prefer the

atomic sites with the largest coordination numbers (CN) while the

smaller osmium atoms prefer the sites of lowest coordination.

The niobium-osmium binary system appears to be

unique in possessing all three of these structures (x, a and p).

At the present time, no other binary system is known to contain

more than two of them.

74
Table 15

Comparison of Site Occupancies for

the X, g and p Phases

CN 12 CN 14 CN 15 CN 16

NbggOSg^

0%Nb

1 OO^cOs

46%Nl3

54'%Os

1 00%Nb

O'^fOs

at 0%Nb

Nbg^Os^Q 100%Os

88%Nb

12%0s

1 00%Nb

0%Os

3* 7%Nb

^N5^^25 93%Os

98%Nb

2tos

* From the present study (See Table 13).

+ From reference 7.

+ From reference 8.

75
A plot of atomic volume versus composition for the

niobium- osmium phases is given in Fig. 39. All three of the

intermediate phases closely follow a linear relationship in this

plot but negative deviations are observed for both terminal solid

solutions.

76
REFERENCES

1. Knapton, A. G. , An X-ray Survey of Certain TransitionMetal Systems for Sigma Phases, Journ. Inst. Metals 87,

pp. 28-32 (1958-1959).

2. Knapton, A. G. , Niobium and Tantalum Alloys, II. Constitution, Structure and Physical Properties,

J. Less-Common Metals 2, p. 113-124 (1960).

3. Dwight, A. E., "Columbium Metallurgy", AIME Metallurgical Society Conferences, Inters cience Publishers, Inc.,

New York, ^0; pp. 383-406 (1961).

4. Taylor, A., Doyle, N. J. and Kagle, B. J., The Constitution

Diagram of the System Molybdenum-Osmium, J. Less-Common

Metals 4, p. 436 (1962).

5. Giessen, B. C, Koch, R. and Grant, N. J., The Niobium

Iridium Constitution Diagram, Trans. AIME 230, p. 1268

(1964).

6. Waterstrat, R. M. The Vanadium- Platinum Constitution

Diagram, Metallurgical Trans. 4j pp. 455-466 (1973).

7. Spooner, F. J. and Wilson, C. G. , Ordering in Binary a Phases, Acta Cryst. 17, pp. 1533-1538 (1964).

8. Van Reuth, E. C. and Waters trat, R. M. , Atomic Ordering

in Binary A15-Type Phases, Acta Cryst. B24, pp. 186-196

(1968).

9. Geller, S., Matthias, B. T. and Goldstein, R. , Some New

Intermetallic Compounds with the "p-Wolfram" Structure,

Journ. Amer. Chem. Soc . 77, pp. 1502-1504 (1955).

77
10. Kasper, J. S., ''Atomic and Magnetic Ordering in

Transition Metal Structures" in Theory of Alloy Phases,

American Society for Metals , Cleveland, Ohio 48A,

pp. 264-278 (1956).

11. Laves, F. , "Factors Governing Crystal Structure" in

Inter metallic Compounds edited by J. H. Westbrook,

John Wiley & Sons, New York, p. 133 (1967),

12. Pearson, W. B. , The Crystal Chemistry and Physics of

Metals and Alloys , Wiley-Interscience Division, John

Wiley & Sons, New York, pp. 55-57 (1972), p. 663 ff.

78
VI. NIOBIUM- PALLADIUM CONSTITUTION DIAGRAM

(Work done at the Massachusetts Institute of Technology

by David Parker and Bill C. Giessen and at the National

Bureau of Standards by Richard M. Waterstrat)

A. Previous Studies

The earliest published information on the constitution

of niobium- palladium alloys is apparently found in a paper by

Greenfield and Beck^^^ who studied seven alloys which were all

annealed at 1000° C. They reported the existence of a sigma phase

and did not observe any ^^^^intermediate phases. A later survey

of this system by Knapton reported an appreciable solubility of

palladium in body-centered cubic niobium (at least 25 at. % Pd)

and the observation of face-centered cubic x-ray powder patterns

for palladium- based alloys containing up to 50% Nb. Knapton was

unable to confirm the presence of a sigma phase in alloys containing 30 to 50 at. % Pd. No sigma structure was detected after

arc-melting or after annealing at 1000 to 1200° C for 5 to 7 days.

In 1961 an equilibrium diagram for the Nb-Pd system

was published by Savitskii, Baron and Khotinskaya^^^. They

reported a sigma phase at 60 wt. % Nb which they chose to

designate as NbgPd. They also reported a compound NbPd^ whose

structure was not determined. The sigma phase was reported as

forming by a peritectic reaction at 1650° C. NbPd^ was described

as being formed at 1700° C from the melt. Extensive mutual solid

solubility was reported for both of the terminal solid solutions.

In 1964 the structures of several Nb-Pd compounds

(4 5^

were identified^ ' and the phases existing in this system were

79
listed as sigma, aNbPd^ NbPdg, aNbPdg and pNbPdg. However,

a somewhat later study reported Nb^-Pdg forming at low

temperatures and Nb2Pd2 (a sigma phase) forming at high temperatures. The compounds NbPd«- and NbPdo were also reported. A

subsequent reinvestigation of the Nb-Pd system^ ' yielded a phase

diagram which differs significantly from the one reported by

Savitskii et al^^^ and identifies the compounds NbgPd, Nb2Pd2,

NbPdg and NbPdg.

B. Materials

The niobium metal used for most of this study was in

the form of -200-mesh powder having a nominal purity of 99. 8%

Some alloys were made using niobium rondelles of 99o6% purity

supplied by the Electro Metallurgical Co. but the use of this material was generally confined to some of the Pd-rich samples. The

Nb powder was obtained from Leico Industries Inc. , New York,

and an analysis of this material after arc- melting showed:

Al, <300 ppm; Ca, <10 ppm; Fe, <1000 ppm; IV^, <100 ppm;

Si, <1000 ppm. Cr, Cu and Ta could not be determined due to

niobium interferences but the supplier has indicated that Ta is the

major impurity. No other elements were detected spectrographically.

Vacuum fusion analysis showed: H, 15 ppm; O, 3866 ppm; C, 53 ppm;

and N, 95 ppm.

The palladium metal used in this study was in the form of

powder or wire, both having a nominal purity of 99.9+%. The

powder was obtained from Engelhard Industries Inc., Newark,

N. J. and the wire from the International Nickel Co. , Inc. Neither

of these materials was submitted to a chemical analysis.

80
C. Alloy Preparation

Weighed amounts of the Nb and Pd powders were

thoroughly mixed and then compressed in a 1/2- inch diameter

cylindrical steel mold. The compressed powder cylinders were

sintered in a high- vacuum furnace prior to arc- melting. Weight

losses after sintering and after arc- melting three or four times

were, in all cases, less than 1%. The arc- melted samples of 5 to

10 grams each were broken or sectioned to check for inhomogeneities but in most cases the samples appeared to be well

melted. A few Nb-rich alloys were prepared using Nb rondelles

and in some of these samples we observed unmelted Nb. Such

alloys were discarded and new alloys were made using powder

mixtures to insure homogeneity. Ultimately all of the samples

were well melted and suitable for preliminary homogenization

treatments. In view of the small weight losses, it was assumed

that the actual alloy compositions did not deviate significantly

from the nominal compositions.

D. Temperature Measurement

Temperatures up to 1500° C were measured with a

platinum versus platinum- 10% rhodium thermocouple. Calibration of furnace temperatures was accomplished by plotting furnace

current versus temperature using a standard platinum-6% rhodium

versus platinum- 30% rhodium thermocouple. This thermocouple

had been calibrated at the National Bureau of Standards. Calibration above 1500° C was accomplished using secondary standards

81
such as the melting points of gold, platinum, rhodium, iridium

and niobium. All temperatures are measured on the International

Practical Temperature Scale of 1968 (IPTS 68). Temperatures in

the vacuum furnace were also measured \vith a Leeds and Northrup

optical pyrometer.

E. Thermal Treatments

All alloys were homogenized by annealing at temperatures near the melting point of each alloy for several hours. The

homogenization treatments were done using the normal precautions

taken during an equilibration treatment and the samples were,

therefore, considered as being equilibrated at the homogenization

temperature. Equilibration annealing of all samples vvas done at

progressively lovver temperatures for each sample.

The annealing treatments above 1100° C were performed in a high- vacuum furnace having a tantalum heating

- fi element and tantalum heat shields. A pressure of betvveen 5 x 10"

- 5 and 2 X 10 mm Hg vvas usually obtained during annealing. Samples

were rapidly cooled after annealing by turning off the furnace

po\ver. The cooling rate was not sufficient to retain all of the

high- temperature structures formed during annealing but information obtained from the partially decomposed samples was sufficient

to permit an understanding of the high- temperature behavior. A

summary of typical annealing treatments is shown in Table 16.

Samples annealed at or below 1100° C were sealed in

quartz tubes under a partial pressure of pure argon and were

water quenched after annealing.

82
Table 16

Typical Equilibration Treatments for

Niobium- Palladium Alloys

Temperature Alloys

rc) Time (Atomic %Pd)

1900 2 hrs

1700 3 hrs

1520 4 hrs

1 400 12 hrs

,

1300 2 days

1200 1 wk

1100 2 wks

1000 3 wks

900 1 month

800 2 months

10, 20

10, 20, 25

25, 50, 52.5, 55, 57.5,60, 64,

70,72. 5, 78, 79,80,81

10, 20, 25, 28, 30, 32, 34, 36, 38,

40,47.5, 50, 52.5,55, 57.5,

60,62.5,65,66.6,67.5,70,

72.5,75, 77,80

10, 20,25, 28,30, 32, 34, 36,

38, 40, 47. 5, 52.5, 55, 57.5,

60,62.5,65,66,6,67.5,70,

72.5, 77,80

10, 20, 25, 28, 30, 32, 40,47. 5,

52.5, 57.5, 62.5, 66.6, 67,

72. 5, 80

25,28,30, 32, 34, 36, 38,40,

52. 5,60,70, 90

25,28,30,32, 34, 36, 38,40,

57. 5,66.6, 72. 5,77

25,28,30, 32, 34, 36,38

25, 36, 38

83
F, Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The procedures used in determining the

(8 solidus temperatures have been described in a previous paper^

High-purity alumina crucibles were found to be satisfactory for

all alloys containing more than 20 at. % Pd but thoria crucibles

were used for the Nb-rich alloys.

2. Liquidus Temperatures

In general, we did not attempt to determine

the liquidus temperatures but they could be estimated with satisfactory accuracy for alloy compositions lying in the vicinity of a

congruent melting maximum or a eutectic minimum. The estimated

liquidus boundaries are shown by dashed lines in Fig. 18.

3. Invariant Reactions

The existence of invariant reactions and the

temperatures at which these reactions occur \vere established

mainly by comparing metallographic and x-ray diffraction data

from alloys annealed above and belosv the transformation temperatures. The phases existing above the transformation temperatures

could not always be retained to room temperature by rapid quenchii^

but metallographic studies were successful in revealing the

structures characteristic of samples annealed either above or

below the transformation temperature. The actual temperature

of the transformation is established by successive annealing

treatments which bracket this temperature with an accuracy comparable to that of the temperature measurement itself.

84
4. Other Phase Boundaries

In previous studies we have relied extensively

on the analysis of individual phase regions in two-phase alloys

using the electron- beam microprobe. Unfortunately, the alloys

in this system were frequently unsuited to this method due to

the presence of unre solvable fine structures in the alloy samples.

These structures originate in the eutectoid decomposition of

phases occurring at high temperatures or by precipitation from

solid solutions. Two separate eutectoid transformations occur

in this system, and we were forced to rely on metaHographic

studies to locate these phase boundaries. The metallographic

studies require more effort and time than the electron probe

measurements, but the results are sufficiently reliable to

establish the location of the phase boundaries to the desired

accuracy.

G. Experimental Results

1. General

The proposed constitution diagram for the

niobium- palladium system is shavn in Fig. 18. The phase boundaries

were established mainly by metallographic and x-ray diffraction

studies. The alloy samples frequently contained phases which had

decomposed from prior high- temperature structures. In all cases,

however, it was possible to obtain a consistent and rational interpretation of these structures.

2. Niobium Terminal Solid Solution - a

The solubility of palladium in body- centered cubic

niobium extends to about 37 at. % Pd at 1520° C. However, there is

85

a significant reduction in solubility at lower temperatures so that

at 800° C the solubility of palladium extends to only about 20 at. % Pd

(See Fig. 19aandl9b). Thus, when alloys containing 25 to 37 at. %Pd

were annealed above 1200° C and cooled rapidly, the cooling rates

employed in this study were apparently insufficient to prevent a

significant precipitation of the intermediate phase 6-NbPd2. This

is shown in Fig. 20a, One can still see the outline of the relatively

large primary grains of the a- phase solid solution but there is

evidence of a precipitate at the grain boundaries and also within

the grains. When alloys containing Pd in excess of the maximum

solubility at high temperatures are quenched and examined metallographically, one observes some additional second phase (sometimes

showing evidence of decomposition) in addition to the phase which

has obviously precipitated (See Fig. 20b, 21a, 21b, 22a). Thus it is

possible to estimate the location of the solvus boundary at high

temperatures by examination of the alloys \vhich exhibit partial

decomposition.

Alloys in this composition range exhibit some

interesting micros tructures when precipitation occurs at

relatively low temperatures, as shown in Fig. 19a and 20b. It is

possible that some of these structures are associated with the formation of a metastable phase. Maldonado and Schubert^^\ for

example, have reported the existence of an orthorhombically

deformed structure based on the body-centered cubic structure

(NbRu-Unterstruktur) after annealing the alloy Nbr^gPdg^ at 650°

C

for 500 hours. This phase was apparently also formed in their

alloy NbggPdg^ after annealing at 700° C for 48 hours and is designated by these authors as Nb2-Pd2(i*) implying that it is a roomtemperature phase.

87
16 ox (a)

75 at. % Nb, 25 at. % Pd equilibrated at 1100°C

for two weeks. Large grains of a Nb containing a

fine network of striations which are possibly

associated with a metastable condition.

w ' 'J

-'^ ^ / .

.4V

/

160X (b)

75 at. % Nb, 25 at. % Pd equilibrated at 1000°C

for three weeks. Oriented precipitation of 6 phase in a matrix of a Nb solid solution.

Fig. 19 - Microstructures of niobium- palladium alloys.

88
160X (a)

72 at. % Nb, 28 at. % Pd equilibrated at 1300°C

for two days and rapidly cooled. Precipitation

of 6 phase occurs due to an insufficient rate of

cooling.

160X (b)

70 at. 9c Nb, 30 at. % Pd equilibrated at 1100°C

for two weeks and quenched. Grain boundaries

contain equilibrated particles of 6 phase. The

interior of each grain contains characteristic

striations formed as a result of an insufficient

quenching rate.

Fig. 20 - Microstructures of niobium- palladium alloys.

89
r

. ... ... ^

'J'

y

i<. . "i

.1

V, -its?

160X (a)

66 at. % Nb, 34 at. % Pd equilibrated at 1300°C

for two days and rapidly cooled. Extensive

precipitation of 6 phase occurs despite the rapid

cooling.

160X (b)

65 at. % Nb, 35 at. % Pd equilibrated at 1100°C

for two weeks and quenched. Matrix is a Nb and

contains globular particles of the § phase.

Fig. 21 - Microstructures of niobium- palladium alloys.

90
160X (a)

64 at. % Nb, 36 at. % Pd equilibrated at 1300°C

for two days and rapidly cooled. Prior structure

consisted of large grains of a Nb with a small

percentage of 5 phase. Extensive precipitation of

5 phase occurs during rapid cooling.

400X (b)

68 at. % Nb, 32 at. % Pd equilibrated for one

month at 900°C. Cellular precipitate of 6 phase

consuming a fine point precipitate.

Fig. 22 - Microstructures of niobium- palladium alloys.

91
However, it appears that one cannot yet rule

out the possibility that this is a metastable phase which is

perhaps formed as a result of cold-isPTOrking the alloy powders

used in the x-ray diffraction studies and then inadequately annealing

these powders. We did not observe this structure in alloys which

were annealed at higher temperatures. In the sample annealed at

^00°C a cellular precipitate was observed which apparently

consumes a prior fine structure during its growth (Fig. 22b).

The occurrence of these precipitate structures

in the a solid solution suggests that interesting mechanical

properties could perhaps be obtained by a suitable heat- treatment

of alloys in this composition range.

3. Intermediate Phase - y'

In the equiatomic region one observes a hightemperature phase, y', which decomposes via a eutectoid reaction

at 1255° C. Alloys quenched from above this temperature contained

the finely-decomposed structure seen in Fig. 23a. We were, thus,

unable to retain the high- temperature phase by rapid quenching.

Alloys annealed below this temperature contained spheroidized

structures produced by completion of the eutectoid reaction (Fig. 23b).

By analogy to the V-Pd system^^^ it is suggested that the y' phase

is probably an extension of the face-centered cubic y solid solution. This could perhaps be checked by high- temperature x-ray

diffraction studies. High- temperature x-ray studies on this

(1) phase have been done by Bowers^ % who reported a cubic pattern,

but due to the lack of data at low angles in this camera he was

unable to determine whether it was, in fact, a face-centered cubic

pattern.

92
4. Intermediate Phase - 6NbPd2

The intermediate phase, 6NbPd2, apparently

forms via a peritectic reaction at about 1610° C. This was established by our solidus data and, particularly, by annealing the alloys

NbggPdgg, NbggPdg^ and Nb^pPdr^Q simultaneously at 1580° C. The

1580° C anneal produced evidence of melting in the microstrue tures

of NbggPdg^ and Nb^gPdgg but no evidence of melting in NbgQPd^^.

The crystal structure of 6NbPd2 has been

determined previously ^''"^ as being an orthorhombic

MoPt2-type structure. This svas confirmed in the present study.

The lattice parameter was not measured since accurate data are

already available^^^\

5. Intermediate Phases - 6 and 9^

The intermediate phase, eNbPd^, has been

reported as possessing an orthorhombic crystal structure based

on ordered close-packed atomic planes vvith a 6-layer stacking

f4 5) (4 5) sequence^ ' It has previously been designated gNbPd^ ' Another structure, e'NbPd^ (previously designated as aNbPd^),

has been reported to be a tetragonal Ti-Al^-type structure^"^' \

These two phases are closely associated in the phase diagram; the

e' phase being slightly more Pd rich. It is not known whether

these two phases are separated by a narrow two-phase region

since we were unable to produce any two- phase alloys in this

region. One of these phases may, in fact, be a high- temperature

(1) phase since Bowers^ ^ has reported evidence of a transformation

at 950° C. We observed that a e phase (NbPd^) alloy transformed

to a e' phase (TiAl^-type structure) after annealing at 850° C.

High- temperature x-ray diffraction studies would help to clarify

this situation.

93
eNbPdg apparently forms at about 1610 ± 10° C

via a congruent order-disorder transformation from the facecentered cubic Y solid solution. This construction is supported by

metallographic evidence and solidus data as summarized in Fig. 18.

6. Palladium Terminal Solid Solution - y

The face- centered cubic palladium solid

solution apparently passes through a congruent melting maximum

at 1630° C and 25 at. % Nb and then extends to about 31 at. % Nb at

high temperatures. It decomposes via a eutectoid reaction at

about 1570°C and 29 at. % Nb (Fig. 24a). However, at temperatures

below 1500° C it appears that the solubility of Nb extends only to

about 18 at. % Nb.

Alloys annealed in the two- phase (e' + y) region

show evidence of a coherent precipitation of 9' in a matrix of the

y phase (Fig. 24b). One might expect to produce some interesting

mechanical properties in alloys containing about 20 at. % Nb.

• Accurate lattice parameters for e and 0' phases

have been reported previously and we have, therefore, acce pted

these values

7. Eutectic: L ^ a + y'

This eutectic reaction was located at 1520 ±

10° C by solidus measurements. A typical eutectic microstructure

is seen in Fig. 23a.

94
160X (a)

52 at. % Nb, 48 at. % Pd in the "as-cast" condition.

Rapidly cooled dendrites (originally y^) containing

a finely decomposed structure of a Nb + § phase. The

interdendritic regions contain a eutectic which was

originally a Nb + The yi, however, has decomposed

eutectoidally both in the primary dendrites and in

the eutectic regions.

42. 5 at. % Nb, 57. 5 at. % Pd. Annealed at lOOCPC

for three weeks. Spheroidized structure produced by

the eu tec to id decomposition of y^ phase.

Fig. 23- Microstructures of niobium- palladium alloys.

95
400X

30 at. % Nb, 70 at. % Pd annealed at 1580 °C

for two hours and rapidly cooled. Relatively

coarse eutectoid structure of 9 + 6 phases produced

by decomposition of the high- temperature y Pd

solid solution.

500X (b)

20 at. 9c Nb, 80 at. f Pd. Annealed at 1357° C

for two days. Coherent precipitate of 9 phase

in a matrix of y Pel,

Fig. 24 - Microstructures of niobium- palladium

alloys.

96
8. Peritectic: L + 6 =^ y'

This reaction was located at 1570 ± 10° C

primarily on the basis of solidus measurements and metallographic

studies as summarized in Fig. 18.

9. Peritectic: L + y' ^ 6

This reaction was established at 1610 ± 10°

C

mainly by careful solidus measurements.

10. Eutectoid: y 6 + e

This reaction was located at about 1575 ± 10°

C

based on metallographic studies of samples cooled rapidly from

various temperatures. A typical eutectoid microstructure is

seen in Fig. 24a.

IL Eutectoid: y' ^ a + 6

This reaction was observed to occur at 1255 ±

5°C based on metallographic studies quenched from above and

below the transformation temperature (Fig. 23a and 23b).

H. Discussion

1. Absence of Sigma Phase

The existence of a binary Nb-Pd sigma phase

(\ 3 4 6 ^ has been reported by various investigators^ >

r > ) ^^^^ somewhat uncertain whether all of these investigators actually

97
observed a sigma phase or whether they simply listed sigma phase

among the phases forming in this system based on the earlier

paper of Greenfield and Beck^^\

(2) Knapton^ ^ apparently made a deliberate attempt

to confirm the observation of sigma phase since he studied alloys

in the appropriate composition range (30 to 50 at. % Pd) and

annealed them in a suitable temperature range (1000 to 1200° C) to

correspond with the observations of Greenfield and Beck^^^.

However, there was no evidence of sigma formation. We made a

similar attempt to reproduce the conditions reported by Greenfield

and Beck by preparing an alloy Nbg^Pd^Q and then using the exact

annealing treatment reported by Greenfield and Beck (72 hrs. at

1000° C). The alloy was two-phase (aNb + SNbPdg) and there was

no evidence of sigma phase. Thus we have confirmed Knapton's

(2) observations^ ' in contradiction to the report of Greenfield and

It is possible that small amounts of impurities

could stabilize the sigma phase and, in this connection, we note that

the Nb metal used by Greenfield and Beck contained a rather large

amount of carbon (0. 5% by weight). It also contained tantalum

(0. 2%), iron (0. 01%), silicon (0. 01%) and titanium (0. 03%). There

were no figures given for the gaseous impurities oxygen, nitrogen

(2) and hydrogen. The niobium used by Knapton^ ^ was apparently

obtained from Johnson, Matthey and Co., Ltd. but no details are

given regarding the purity of this material. Our own niobium had

a much lower carbon content (53 ppm) than the material used by

Greenfield and Beck. At any rate, it appears that sigma phase did

not form in any of our alloys after annealing at progressively

lower temperatures in 100° increments from 1400° C to 800° C.

98
This is particularly surprising since In the Ta-Pd system the sigma

phase is stable up to approximately 2550°C^^^\ The V-Pd system^^^

contains an A15-type phase in this range of Pd content but no A15-

type phase was detected in our Nb-Pd alloys even though we

annealed the alloy Nb^Pd and others at various temperatures down

to 800° C. It, therefore, appears that the Nb-Pd system is unique

among these three systems in not forming either a sigma or an

A15-type phase. Instead it appears that there is a greater solubility of Pd in the body- centered cubic aNb solid solution than is the

case with Pd in the aTa solid solution. This observation may be of

some practical importance in developing useful Nb or Pd alloys by

suppressing the formation of the brittle compounds.

2. Alloying Behavior

Exploration of ternary systems such as Ta-Nb-Pd

may suggest some reason for the remarkable behavior of Nb in

suppressing these phases. We have prepared several alloys in

the Ta-Nb-Pd system and annealed them for three weeks at 1100° C.

We observed the sigma phase as a major constituent of ternary

alloys with Ta contents ranging from 15 to 20 at. %. The alloy

Ta^^gNbggPdgQ was essentially all sigma despite its low Ta content.

We also prepared several alloys in the V-Nb-Pd system and observed Sigma phase as a major constituent for V contents ranging

from 15 to 20 at. %.

Thus it appears that small amounts (< 15 at. %)

of either V or Ta, when added to Nb-Pd alloys, will stabilize the

brittle sigma phase. Previous reports of a sigma phase forming

in the Nb-Pd binary system may, therefore, be in error; probably

due to the effects of a slight contamination of the experimental

alloys. Crystallographic data for Nb-Pd alloys are summarized

in Tables 17 and 18.

99
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101
REFERENCES

1. Greenfield, P. and Beck, P. A., Intermediate Phases in

Binary Systems of Certain Transition Elements, Trans.

AIME 206, pp. 265-276 (1956).

2. Knapton, A. G. , Niobium and Tantalum Alloys. 11 Constitution,

Structure and Physical Properties, J. Less-Common

Metals, 2, pp. 113-124 (1960).

3. Savitskii, E. M., Baron, V. V. and Khotinskaya, A. N.

,

Equilibrium Diagram of the Niobium- Palladium System , Russian Journal of Inorganic Chemistry, 6, pp. 1316-1317 (1961).

4. Giessen, W. C, Parker, D, and Grant, N. J., The Intermediate Phases Ab, Abg and Ab^ in the Nb-Pd and Nb-Pt

SystemSo Journ. Metals (abstract) pp. 92-93 Jan. (1964).

5. Giessen, B. C. and Grant, N. J., New Intermediate Phases

in Systems of Nb or Ta with Rh, Ir, Pd or Pt, Acta Cryst.

,

VI pp. 615-616 (1964).

6. Maldonado, A. and Schubert, K. Structure Investigation in

5 10

the T -T Homologous and Quasihomologous Alloy Systems,

Z. Met. 55^ pp. 619-626 (1964).

7. Bowers, I., personal communication to B. C. Giessen (1968).

8. Waterstrat, R. M., The Vanadium- Platinum Constitution

Diagram, Met. Trans. 4, pp. 455-466 (1973).

9. Koster, W. and Haehl, W. D. , The Binary System PalladiumVanadium, Z. Met. 49 , pp. 647-649 (1958).

10. Schubert, K. , Frank, K. , Gohle, R., Maldonado, A.,

Meissner, H. G. , Raman, A., and Rossteutsc her, W.

,

Some Structure Data for Metallic Phases, Naturwiss.

50, p. 41 (1963).

102
11. Giessen, B. C. and Grant, N. J., The Crystal Structure

of VNig, VPdg, VPtg and Related ABg Phases. J.

Less-Common Metals 8, pp. 114-119 (1965).

12. Dwight, A. E., " Columbium Metallurgy", AIME Metallurgical Society Conferences 10, pp. 383-406 (1961).

Interscience Publishers, Inc. , New York.

13. Waterstrat, R. M., Koch, R. and Giessen, B. C, The

Tantalum- Palladium Constitution Diagram, To be

published (1973).

103
VII. NIOBIUM- PLATINUM CONSTITUTION DIAGRAM

(Work done at the Massachusetts Institute of Technology

by Bill C. Giessen and at the National Bureau of Standards

by Richard M. Waterstrat and Richard C. Manuszewski)

A, , Previous Studies

Identification of intermediate phases in the niobiumplatinum system was begun by Greenfield and Beck^^^ who

identified a sigma phase and a phase having an A15 (p'-W)-type

structure. Knapton^^^ later reported that the liquidus in this

system passes through a maximum near the composition NbPt^.

In 1961 a phase diagram for the platinum- niobium system was

presented by Kimura and Ito^^^ who confirmed the presence of the

Sigma and 3-W type phases and also reported an inter metallic

compound having the approximate composition NbPt^ but the

crystal structure of this compound was not determined. There

was some evidence of a compound at the equiatomic (NbPt)

composition.

(A 5^

In 1964 Giessen et ar ' ' reported that the equiatomic

compound, NbPt, possesses a B19 or AuCd-type structure and that

another phase occurs near this composition which is only stable at

high temperatures (1680-1770°C). They also reported ^^"^^ a

compound, NbPtg, having an orthorhombic structure isomorphous

with TaPd2(MoPt2-type) and two compounds at the approximate

composition NbPt^. aNbPt^ was reported as having an orthorhombic

DO^ or TiCug-type structure and pNbPt^ was identified as a monoclinic phase ^vith a 12-layer stacking sequence isomorphous vvith

pTaPtg^^"^^. It has been suggested^^^ that aNbPtg is probably

stabilized with respect to pNbPt^ by very small amounts of inter104
stitial impurities. However, Maldonado and Schubert^ ' indicate

that pNbPtg may have a somevvhat higher platinum content

(NbPt«4.) than aNbPtq. All of these phases, previously reported,

have been confirmed by Maldonado and Schubert. ^

'

B. Materials

The niobium metal used in this study was in the form of

pellets or rondelles which were obtained from both the Electro

Metallurgical Co. and from the Gallard Schlesinger Chemical

Manufacturing Corp. , New York. The nominal minimum purity of

this material was 99. 6% with the principal impurity being

tantalum. A spectrographic analysis showed: Ag, < 10 ppm;

Al, < 300 ppm; Ca, < 100 ppm; Fe, < 1000 ppm; Mg, < 100 ppm;

Mn, < 10 ppm; Mn, < 10 ppm; Ni, < 10 ppm; Si, < 300 ppm. Cr,

Cu and Ta could not be determined due to niobium interferences but

all other impurities were less than 10 ppm.

The platinum metal used in this study was in the form

of 1/32-inch thick sheet from an unknown source. It had a nominal

purity of 99. 9% and spectrographic analysis revealed: Fe, < 10 ppm;

Cu, < 100 ppm; Mg, < 100 ppm; Si, < 100 ppm; Pd, <1000 ppm;

Ru, < 100 ppm; all other impurities, < 10 ppm.

C. Alloy Preparation

Carefully weighed quantities of Nb and Pt were arcmelted at least four times using a sufficient arc current to insure

that the alloy button was thoroughly melted. Weight losses during

melting were in all cases less than 1% except for the alloy, NbggPt^,

which had a weight loss of 1. 5%. The actual alloy compositions were

105
therefore assumed to be reasonably close to the intended nominal

compositions. Each alloy sample contained 20 grams of material

and was sectioned or broken for visual inspection prior to homogenization treatments. In all cases the samples were ^vell melted and

no problems were encountered in producing them,

D. Temperature Measurement

Temperatures up to 1500° C were measured with a

platinum versus platinum-10% rhodium thermocouple. Calibration

of furnace temperatures was accomplished by plotting furnace current

versus temperature using a standard platinum-6% rhodium versus

platinum- 30% rhodium thermocouple. This thermocouple had been

calibrated at the National Bureau of Standards. Calibration above

1500° C was accomplished using secondary standards such as the

melting points of gold, platinum, rhodium, iridium and niobium.

All temperatures are measured on the International Practical

Temperature Scale of 1968 (IPTS 68). Temperatures in the vacuum

furnace were also measured with a Leeds and Northrup optical

pyrometer.

E. Thermal Treatments

All alloys were first homogenized by annealing them at

temperatures slightly below their melting points for several hours.

The samples could be considered as being equilibrated at these

temperatures and were rapidly cooled by turning off the furnace

power. Further, equilibration treatments were then conducted at

progressively lower temperatures.

106
The annealing treatments above 1100° C were done in a

high vacuum (-5 x 10" mm Hg) using a Ta-strip heating element

and Ta heat shields. There was no evidence of unusual weight

losses or contamination of the samples. Nb foil was sometimes

used to wrap or hold the samples but in most cases Ta foil seemed

adequate for this purpose. Alloys annealed at or below 1100°

C

were sealed in quartz tubes under a partial pressure of argon.

These tubes were water-quenched following their annealing treatments. A summary of typical annealing treatments is shewn in

Table 1 9.

F. Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The procedures used in determining the solidus

temperatures have been described in a previous paper^^^\ Highpurity alumina crucibles were found to be satisfactory for all

alloys containing more than 20 at. % Pt but thoria crucibles vvere used for the Nb-rich alloys.

2. Liquidus Temperatures

In general, we did not attempt to determine the

liquidus temperatures but they could be estimated with satisfactory accuracy for alloy compositions lying in the vicinity of a

congruent melting maximum or a eutectic minimum. The estimated

liquidus boundaries are shown by dashed lines in Fig. 25.

3. Invariant Reactions

The existence of invariant reactions \vas established by solidus measurements and metallographic studies on

107
Table 19

Typical Equilibration Treatments for

Niobium- Platinum Alloys

Temperature Time Alloys

(°C ) (Atomic % Pt)

1900 3 hrs 16,18, 20

1 700 6 hrs 5, 10, 16, 18, 22, 26, 30, 50,

52.5, 55, 57.5,60, 62i 5,

66.6

1600 12 hrs 45, 50, 51,52.5,55,57.5,

'

60,62.5,67.5, 70

1500 Iday 5,10,18,22, 26, 30,32, 34,36,38,

45,50,52.5,55,57.5,60,62.5,

72, 73, 77. 5, 80,82.5,85,

90,95

1300 1 wk 45, 50,52.5, 55, 57.5,60,

62.5,73,77.5

11 50 3 wks 5, 10, 16, 18, 20, 22, 26, 30, 32,

34, 36,38, 50,70,77.5,80,

82. 5,85,90, 95

1000 1 month 50

108
rapidly-cooled samples. We obtained no evidence supporting the

existence of any solid-state reactions in this system other than

precipitation from the saturated terminal solid solutions.

4. Other Phase Boundaries

Phase boundaries were generally determined by

metallographic and x-ray diffraction studies but in many cases we

were also able to make use of the electron-beam microprobe to

obtain individual chemical analyses of the phases coexisting in twophase alloys equilibrated at various temperatures. This was of

great help in locating the phase boundaries except for samples in

which precipitation had occurred in one of the coexisting phases.

In such cases, the use of metallography was our only recourse. It

is estimated that the phase boundaries shown in Fig.25 are accurate

±1 at. %.

G. Experimental Results

1. General

The proposed constitution diagram is shown in

Fig.25. A discussion of the individual phase regions follows. All

alloys were examined either by x-ray diffraction or metallographically or both. In addition, x-ray microprobe data were used to

establish the location of certain phase boundaries. The accuracy

of these boundaries is believed to be about ±1 at. % in composition.

2. Niobium Terminal Solid Solution - a

Platinum is soluble in the body-centered cubic

niobium solid solution to the extent of about 14 at. % Pt at 2040° C

109
Jo'3anivd3dW3i
but the solubility diminishes to about 5 at. % Pt at lower temperatures and precipitation of a second phase is frequently seen in

alloys cooled slowly enough to permit such effects (Fig.26a). The

a-phase lattice parameters are affected by Pt in solid solution

which gradually decreases the lattice constant of pure Nb.

3. Intermediate Phase -

p

This phase was apparently first discovered by

Geller, Matthias and Golds tein^^^ who also reported on the degree

of atomic ordering. Their discovery was confirmed by Greenfield

(I) (3 4 7^ and Beck^ ^ and later by other workers'' ' ' \ The structure is a

cubic A15 (p-W) type and the phase extends from about 19 at. % Pt

at 2040° C to about 28 at. % Pt at 1800° C. It apparently forms by a

peritectic reaction at 2040° C.

4. Intermediate Phase - a

This phase was first reported by Greenfield and

Beck^^^ who identified its tetragonal crystal structure as that of a

new Sigma phase. It exists over a composition range extending

from about 31 to 38 at. % Pt at high temperatures but its range

diminishes at lower temperatures (See Fig. 26b). Its presence in

13 4 7) the Nb-Pt system has been confirmed by various investigators^ ' ' \

It apparently forms by a peritectic reaction at 1800° C.

5. Intermediate Phase -

Q

This phase was first identified by Giessen

(4 5) et ar ' ' as having an orthorhombic AuCd-type structure. It is

apparently stable from its solidus point at about 1750° C down to

below 1000° C. It coexists with the sigma and 6 phases (Fig. 2 7a

and 27b).

Ill
400X (a)

82 at. % Nb, 18 at. % Pt. Annealed at 1150 C for

three weeks. Beta phase matrix containing a Nb

globules in which further precipitation of beta has

occurred.

160X (b)

70 at. % Nb, 30 at. % Pt. Annealed at 1500 C

for one day. Beta phase in a matrix of sigma phase.

Fig. 26 - Micros true tures of niobium- platinum alloys.

112
160X (a)

62 at. % Nb, 38 at. % Pt. Annealed at 1500 C for

one day. c; phase globules in a matrix of sigma phase.

42. 5 at. % Nb, 57. 5 at. % Pt. Annealed at 1000 C for

one month. 6 phase globules in a matrix of phase.

Fig. 27 - Microstructures of niobium- platinum alloys.

113
An analogous phase having this same crystal structure occurs in

the V-Pt binary system^^^^ where it is also observed at the equiatomic composition but is stable only up to about 1500° C.

6. Intermediate Phase - y'

The existence of this high- temperature phase has

been deduced^ ' from solidus measurements and metallo graphic

studies. It apparently forms by a peritectic reaction at 1790° C and

decomposes below 1680° C to produce the fine structure shown in

Fig. 28a. By analogy to the V-Pt system^ ' we would expect this

phase to be essentially an extension of the face- centered cubic y

solid solution and have therefore labelled it y' . Confirmation of

its crystal structure, however, must await a high- temperature x-ray

diffraction study.

7. Intermediate Phase - 6

This phase forms over a narrow composition range

around the composition NbPto. It possesses an orthorhombic crystal

M 7

structure of the MoPtg-type ^ > ^ ' . It probably forms by a peritectic reaction at about 1990° C. Accurate lattice parameter values have

been reported previously and were accepted in this study^^'

8. Intermediate Phases - e and e'

Two different crystal structures have been observed

at the composition NbPto. The e' phase (formerly designated

aNbPt^)^ ' ' ' has an orthorhombic DO„ or TiCu^-type structure. The

e phase (formerly designated BNbPt«) ^ ' ' has a monoclinic structure

(4-7^ with a 12- layer stacking sequence and is isomorphous with p-TaPtq^ '

(<7)

«^ Maldonado and Schubert^ ' have stated that the TiCu«-type structure is

114
stable beloNV 1000° C and they designate the TaPt^-type structure as

having a Pt-content which is somewhat higher than the stoichiometric

(Abg) composition. High- temperature x-ray diffractometry may

help to clarify the structural behavior. The 9 phase melts congruently at about 2030° C. Accurate lattice parameters for the e and

phases have been reported previously and were accepted in this

study^ ' '. 9 phase coexists with 6 phase (Fig. 28b) and vPt (Fig. 29a).

9. Platinum Terminal Solid Solution - y

Face-centered cubic platinum can dissolve up to

20 at. % Nb at 2000° C with some small reduction in solubility

occurring at lower temperatures. Additions of Nb produce a sharp

rise in the solidus and liquidus of the Pt solid solution and a gradual • increase in its lattice parameter. This may perhaps be accompanied

by some interesting mechanical properties.

10. Peritectic: a + 3

The existence of this peritectic reaction was

established by solidus measurement and by observation of a typical

peritectic microstructure. The reaction temperature is 2040 ± 10° C.

11. Peritectic: p + L ^ ct

This peritectic reaction was located at 1800 ±

10° C mainly by solidus measurements in this composition region.

12. Eutectic: L ^ a + C

The eutectic reaction at 1700 ± 10° C was determined

by solidus measurement and by observation of typical eutectic microstructures (Fig. 2 9b).

115
400X (a)

45 at. % Nb, 55 at. % Pt. Annealed at 1600'='C for

twelve hours, c phase plus 6 phase fine structure

probably produced by decomposition of prior -y* phase

160X (b)

30 at. % Nb, 70 at. % Pt. Annealed at 1800°C for

four hours. 9 phase globules in a matrix of 6 phase.

Fig. 28 - Microstructures of niobium- platinum alloys

116
s?.

160X (a)

20 at. % Nb, 80 at. % Pt. Annealed at 1500°C for

one day. Coherent precipitate of e phase in y Pt.

400X (b)

55 at. % Nb, 45 at. % Pt. ''As-cast" fine eutcctic

of Sigma plus ^ phases.

Fig. 2 9 - Microstructures of niobium-platinum alloys.

117
13. Peritectic: L + ^ C

This reaction occurs at about 1750 ± 10°

C

on the basis of solidus measurements.

14. Peritectic: L + 6 ^

Solidus data suggest that this reaction occurs at

about 1790 ± 10° C.

15. Peritectic: L + 9 ^ 6

Solidus measurements in this region suggest a

peritectic reaction occurs at about 1990 ± 10° C.

16. Congruent melting maximum: L ^ 9

It is estimated on the basis of solidus measurements that the 9 phase melts congruently at 2030 ± 10° C.

17. Peritectic: e + L =^ y

This peritectic reaction apparently occurs at about

2000 ± 10° C on the basis of solidus measurements.

H. Discussion

The proposed constitution diagram of the Nb-Pt

system shares some similarities with the V-Pt^^^^ and the Ta-Pt^'^^

systems. The crystal structures of the ^ Q, b and 9 phases are

all reproduced in at least two and sometimes in all three of these

systems (See Tables 20 and 21). There is a notable tendency toward

an increased stability of the p phase as one proceeds from the Ta-Pt

118
to the Nb-Pt to the V-Pt systems. This is accompanied by a

decreasing stability for the sigma phase. Each of these three

systems also exhibits ordered close-packed structures at the

AB2 and ABg stoichiometries where, in these cases, "B" is

platinum. It would, therefore, appear that similar electronic

conditions exist in these three related systems which may be

attributed to periodic table relationships of the constituent elements.

Further study of physical properties in these three systems may

yield some insight into the reasons for the simularities.

119
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121
RE FERENCES

' 1. Greenfield, P. and Beck, P. A., Intermediate Phases in

Binary Systems of Certain Transition Elements, Trans.

AIME 206, p, 265 (1956).

2. Knapton, A. G. , Niobium and Tantalum Alloys. II,

Constitution, Structure and Physical Properties, J.

Less-Common Metals 2, pp. 113-124 (1960).

3. Kimura, H. and Ito, A. , The Platinum- Niobium System,

Trans. Nat'l. Res. Inst, for Metals 3, pp. 27-34 (1961).

4. Giessen, B. C. , Parker, D. and Grant, N. J., The Intermediate Phases Ab, Abg and ABg in the Nb-Pd and Nb-Pt

Systems, Journ. Metals (abstract), pp. 92-93 (1964)o

5. Giessen, B. C. and Grant, N. J. , New Intermediate Phases

in Systems of Nb or Ta with Rh, Ir, Pd or Pt, Acta Cryst.

pp. 615-616 (1964).

6. Giessen, B. C, Kane, R. H. and Grant, N. J., On the

Constitution Diagram Ta-Pt Between 50-100 At. Pet. Pt,

Trans. Met. Soc. AIME 233, pp. 855-864 (1965).

7. Maldonado, A. and Schubert, K. Structure Investigation in

5 10

the T -T Homologous and Quasihomologous Alloy Systems,

Z. Met. 55, pp. 619-626 (1964).

8. Geller, S. ; Matthias, B. T. ; and Goldstein, R. , Some New

Intermetallic Compounds with "p-Wolfram" Structure, Journ.

Amer. Chem. Soc . 77, p. 210 (1955).

9. Giessen, B. C. and Grant, N. J. , The Crystal Structure of

VNig, VPdg, VPtg and Related ABg Phases. J. LessCommon Metals 8^ pp. 114-119 (1965).

10. Waterstrat, R. M. , The Vanadium- Platinum Constitution

Diagram, Met. Trans . 4^ pp. 455-466 (1973).

122
Vni. TANTALUM- PALLADIUM CONSTITUTION DIAGRAM

(Work done at the National Bureau of Standards by

Richard M. Waterstrat and at the Massachusetts Institute

of Technology by Ros marie Koc h and Bill C . Giessen)

A. Previous Studies

The earliest work on the tantalum- palladium system

is apparently that of Greenfield and Beck^^^ who prepared only

three alloys, annealed them at 1000° C and concluded that no intermediate phases exist at this temperature. Later work by Darby

(2)

et ar however, established the existence of several intermediate

phases. A sigma phase was discovered and a phase having a TiAl^-

type structure was reported; the latter near the composition TaPd^.

A body-centered tetragonal BU-type structure was found at the

equiatomic composition and an unidentified phase (X) was observed

at the composition TaPdo. Several of these phases were identified

(3) in a subsequent study by Kane, Giessen and Grant^ \ The equiatomic phase was confirmed as havir^ a BU or yTiCu-type structure

and the X-phase (TaPd^) was identified as having a Mo- Pt^- type

(3 4) structure^ ' ' isomorphous with NbPt^. The presence of a sigma

(3) phase and a TiAl^-type structure was also confirmed^ A further

confirmation of all these phases was provided by Maldonado and

(5) Schubert^ Finally, a phase diagram for the palladium- tantalum

system was presented by Savitskii et al^^^ but it is based on a rather

small amount of experimental data and the phase boundaries are

apparently not precisely defined. Nevertheless, intermetallic

compounds were observed at the compositions mentioned previously

and their presence is thus again confirmed.

However, it appears that more work is needed to

delineate the various phase regions in this diagram and to insure

123
the reliability of the final result. We have, therefore, undertaken

a complete reinvestigation of the tantalum- palladium system with

a view toward obtaining a more reliable phase diagram.

B. Materials

The tantalum metal used in this study was either in the

form of rod having a purity of 99.95% or in the form of powder

having a purity of 99. 99%. The rod \vas obtained from National

Research Corp. and the powder was obtained from Leico Industries

Inc. The impurities in these materials consisted of O, C, N, Si,

Nb, Mo and W but none of these impurities were present in excess

of 100 ppm.

The palladium metal used in this study was in the form

of powder or wire, both having a nominal purity of 99.9+%. The

powder was obtained from Engelhard Industries Inc. and the wire

from the International Nickel Co. Inc. Neither of these materials

was submitted to chemical analysis.

C. Alloy Preparation

Alloys rich in Pd could be easily prepared using Ta rod

and Pd wire as starting materials but when these materials were

used for Ta-rich alloys there were significant metal losses during

melting and in many cases we observed large pieces of unmelted Ta

in the alloy after it had been broken or sectioned. Eventually we

changed our procedures for the Ta-rich alloys and used powder

metallurgical methods to prepare these alloys for arc- melting.

Weighed amounts of Ta and Pd powders were thoroughly mixed and

124
then compressed in a l/2-inch diameter cylindrical steel mold.

This was followed by a short sintering at about 1200°C in a high

vacuum and then by arc- melting of the pre-sintered compacts. The

Ta-rich powder alloys were each of about 10 grams and appeared to

be well- melted. The Pd-rich alloys contained about 5 grams of

material. Weight losses were in all cases less than 1% during

melting.

D. Temperature Measurement

Temperatures up to 1500° C were measured with a

platinum versus platinum- 10% rhodium thermocouple. Calibration of furnace temperatures was accomplished by plotting

furnace current versus temperature using a standard platinum- 6%

rhodium versus platinum- 30% rhodium thermocouple. This

thermocouple had been calibrated at the National Bureau of

Standards. Calibration above 1500° C was accomplished using

secondary standards such as the melting points of gold, platinum,

rhodium, iridium and niobium. All temperatures are measured on

the International Practical Temperature Scale of 1968 (IPTS 68).

Temperatures in the vacuum furnace were also measured \vith a Leeds and Northrup optical pyrometer.

E. Thermal Treatments

All alloys were homogenized by annealing at tempera - tures near the melting point of each alloy for several hours. The

homogenization treatments were done using the normal precautions

taken during an equilibration treatment and the samples were,

therefore, considered as being equilibrated at the homogenization

temperature. Equilibration annealing of all samples vvas done at

progressively lo\ver temperatures for each sample.

125
The annealing treatments above 1100° C were performed in a high- vacuum furnace having a tantalum heating element

and tantalum heat shields. A pressure of between 5 x 10" and

- 5

2 X 10 mm Hg was usually obtained during annealing. Samples

were rapidly cooled after annealing by turning off the furnace power.

The cooling rate was not sufficient to retain all of the hightemperature structures formed during annealing but information

obtained from the partially decomposed samples was sufficient to

permit an understanding of the high- temperature behavior. A

summary of typical annealing treatments is shown in Table 22.

Samples annealed at or below 1100° C were sealed in

quartz tubes under a partial pressure of pure argon and were water

quenched after annealing.

Table 22

Typical Equilibration Treatments for Tantalum- Palladium Alloys

Temperature

rc)

Time Alloys

(Atomic % Pd)

2300 3 hrs 20,25

2000 6 hrs 13, 20, 25

1680 1 day 13, 20, 25, 30, 40, 48, 50, 53, 55,

57,60,65, 66.6,67, 70, 74, 75,

76,77.5, 80,82, 84,90

1350 1 wk 13, 20, 25, 30, 40, 47. 5, 49, 50,

51, 52, 55, 56, 57,60,65,66.6,

67, 70,74,75

1150 1 month 13, 25,30,47.5, 49,50, 51, 52,

55, 65, 67, 74, 76, 80, 82, 84, 90

1000 1 month 25,40, 50, 51, 55, 60,66.6, 67

70,75,80,82,84,90

126
F. Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The procedures used in determining the solidus

(1) temperatures have been described in a previous paper ^ Highpurity alumina crucibles were found to be satisfactory for all alloys

containing more than 30 at. % Pd but thoria crucibles were used

for the Nb-rich alloys.

2. Liquidus Temperatures

In general, we did not attempt to determine the

liquidus temperatures but they could be estimated with satisfactory

accuracy for alloy compositions lying in the vicinity of a congruent

melting maximum or a eutectic minimum. The estimated liquidus

boundaries are shown by dashed lines in Fig. 30.

3. Invariant Reactions

The existence of invariant reactions and the

temperatures at which these reactions occur vvere established

mainly by comparing metallographic and x-ray diffraction data

from alloys annealed above and below the transformation temperatures. The phases existing above the transformation temperatures

could not always be retained to room temperature by rapid quenching

but metallographic studies were successful in revealing the

structures characteristic of samples annealed either above or

below the transformation temperature. The actual temperature

of the transformation is established by successive annealing

treatments which bracket this temperature with an accuracy comparable to that of the temperature measurement itself.

127
4. Other Phase Boundaries

In most cases it was possible to supplement

our metallographic studies with electron microprobe data. The

only limitation on the use of the probe in this system was for

alloys containing decomposed y' phase. The decomposition of y'

produces a fine structure which is unresolvable to the electron

beam in the microprobe. In some cases where samples contained

unmelted Ta or when the surface of the sample had been depleted

of Pd by annealing in a vacuum, it was decided to treat the samples

as if they were diffusion couples and analyze the phases along a

path perpendicular to the interface between the phases. This

assumes that equilibrium conditions prevail along this path and we

were unable to test this assumption. The data obtained, however, , are consistent with metallographic data in the same composition

region.

G. Experimental Results

1. General

The proposed constitution diagram for the tantalumpalladium system is shown in Fig. 30. It differs significantly from

the diagram reported previously^^\ Individual phase regions are

discussed below.

2. Tantalum Terminal Solid Solution - g

Palladium is soluble in the body- centered cubic

tantalum solid solution to the extent of about 19 at. % Pd at 2550°C

but there is a decrease in solubility at lower temperatures to about

10 at. % Pd at 1200°C (See Fig. 31a). The solvus boundary was

128

located mainly by quantitative analysis of the a-phase in twophase alloys using the electron beam microprobe. The alloys

were equilibrated at various temperatures prior to each analysis.

In some cases it was possible to traverse a reaction zone between

an unmelted piece of pure Ta and a partly melted Pd-rich region.

This, in effect, is a diffusion couple which may or may not correspond to equilibrium conditions. In this case, however, the data

were consistent with data obtained from conventional metallographic

specimens in this composition range. The aTa solid solution

coexists with the Sigma phase (See Fig. 31b).

3. Intermediate Phase - a

This phase was originally discovered by Darby

et al^^^ and later confirmed by other authors ^' ^\ It possesses

a tetragonal crystal structure and it apparently forms by a peritectic

reaction at 2550°C. It is rather remarkable that the sigma phase is

stable to such an elevated temperature in this system while, in the

Nb-Pd system, it does not appear at all, at least above 800°C.

4. Intermediate Phase -

The existence of this phase was established

mainly by the observation of finely decomposed microstructures

in alloys quenched from above 1550° C. This microstructure

(Fig. 32a) is characteristic of a eutectoid decomposition. By analogy

to the V-Pd system^^^ and the V-Pt system^'^^ one might expect

either that the y' phase is essentially an extension of the facecentered cubic Y solid solution or, alternatively, that it possesses

a AuCd-type structure. Clarification of this question must await

further studies.

130
400X (a)

87 at. % Ta, 13 at. % Pd. -Annealed at 1150°C for

one month. Widmanstatten precipitate of sigma

phase in a phase.

80 at. % Ta, 20 at. % Pd. Annealed at 1500°C for

one day. Particles of a phase in a matrix of brittle

Sigma phase.

Fig. 31 - Microstructures of tantallum- palladium alloys

131
5. Intermediate Phase -

The existence of a body- centered tetragonal BU

(2)

or yTiCu-type structure was discovered by Darby et ar ' and

confirmed by other investigators^^' ^' ^\ It forms by a peritectoid

reaction at about 1410° C (See Fig. 32b and 33a) and has a rather

narrow composition range near the equiatomic composition. Its

lattice parameter has already been accurately measured^^^ and

this value was accepted in the present study. It coexists vvith 6

phase (Fig.33a) and sigma phase (Fig. 33b and 34a).

6. Intermediate Phase - 6

The 6 phase TaPdg exists over a narrow composition range near the stoichiometric composition and possesses an

orthorhombic structure of the MoPtg-type^^' ^\ It melts congruently

at about 1800° C. An accurate lattice parameter has already been

measured for this phase which we have adopted.

7. Intermediate Phase -

The e phase, TaPdg, has a tetragonal structure

of the TiAl^-type and exists within a narrow composition range at

the ideal stoichiometry. It melts congruently at about 1780° C.

The published lattice parameter values were adopted in this

study.

8. Palladium Terminal Solid Solution - y

The face- centered cubic palladium solid solution

dissolves up to about 22 at. % Ta at 1730° C but the solubility

132
50 at. % Ta, 50 at. % Pd. Annealed at 1575° C for

six hours. Fine eutectoid structure produced by

decomposition of prior phase.

50 at. %. Ta, 50 at. % Pd. Annealed at 142 5°C for

one day. Islands of sigma phase in a matrix of

6 phase.

Fig. 32 - Microstructures of tantallum-palladium alloys.

133
400X (a)

50 at. % Ta, 50 at. % Pd. Annealed at 1150PC for one

month. Particles of 6 phase in a large- grained

matrix of ti phase.

70 at. % Ta, 30 at. % Pd. Annealed at 1150° C for

one month. Islands of ri phase in a matrix of sigma

phase.

Fig. 33- Microstructures of tantallum- palladium alloys.

134
400X (a)

52. 5 at. % Ta, 47. 5 at. % Pd. Annealed at 1150° C

for one month. Particles of sigma phase in a matrix

of Ti phase.

400X (b)

'^O at. % Ta, 40 at. % Pd. Partially melted at 1700°C.

Eutectic structure produced by partial melting is

composed of light sigma phase and partially decomposed

phase.

Fig. 34 - Microstructures of tantallum-palladium alloys.

135
diminishes at lower temperatures until at 1000° C it is probably

not more than 18 at. % Ta. It is possible that the mechanical

properties of alloys in this composition range would be of some

interest.

9. Peritectic: a + L ^ a

The existence of a peritectic reaction at 2550°

C

was established mainly by a solidus measurement on the alloy

TagPd followed by metallographic study of the solidus sample.

Solidus measurements on alloys with higher Pd contents indicated

that a very sharp decrease in the solidus boundary occurs with

increasing Pd content. Traces of liquid were found, for example,

in the alloy Tar^gPd^Q after annealing at only 1700° C.

10. Eutectic: L g + y

^

Solidus measurements and metallographic studies

(Fig. 34b) established a eutectic reaction at about 46 at. % Pd and

1700 ± 10°C.

11. Peritectic: L + 6

This reaction was indicated mainly by solidus

measurements as occurring at about 1720 ± 10° C.

12. Eutectic: L ^ 6 + 9

Solidus measurements and metallographic

studies (Fig. 35a) located this reaction at about 72 at. % Pd and

1720 ± 10° C.

136
400X (a)

30 at. % Ta, 70 at. % Pd. (Melted) Eutectic

structure of 5 plus 3 phases with primary grains

of 6 phase.

400X (b)

60 at. % Ta, 40 at. 9f Pd. Annealed at 1575 C for

one day. Sigma phase (light) plus finely decomposed

V* phase (dark).

Fig. 35 - Microstructures of tantallum- palladium

alloys.

137
400X (a)

50 at. % Ta, 50 at. % Pd. Annealed at 147 5°C for

one day. Partially spheroidized structure of

decomposed y*.

Fig. 36- Microstructure of a tantallum- palladium alloy.

138
13. Peritectic: e + L ^ y

Solidus measurements indicated a continuous

rise in the melting range of the Pd terminal solid solution with

increasing Ta content terminating in a peritectic reaction at

about 1730 ± 10°C.

14. Eutectoid: y

^ ^ g + 6

This eutectoid reaction was established mainly

by metallographic studies which showed a typical fine eutectoid

structure in alloys quenched from above 1550° C (Fig. 32a, 34b and

35b).

15. Peritectoid: a + 5 =^ n

The existence of this peritectoid reaction was

established mainly by metallographic studies of alloys quenched

from above and below the transformation temperature (See Fig. 32b

and 33a).

H. Discussion

The Ta-Pd phase diagram possesses some unique

features when compared with the systems Nb-Pd and Ta-Pt, for

which one might expect to observe analogous behavior. The Ta-Pd

system possesses a very stable sigma phase in sharp contrast to

the Nb-Pd system where no sigma phase has been observed, at

least above 800° C^^^. The Ta-Pd system also contains a Bll-type

phase at the equiatomic composition in contrast to the Ta-Pt system

where no Bll-type phase is observed, at least above 1000°C^^^^o

139
However, all three of these systems form intermediate phases

with the ABg and AB^ stoic hio metry. Measurements of physical

properties in these three systems could be helpful in revealing

the reasons for the significant differences in their alloying behavior.

140
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REFERENCES

1. Greenfield, P. and Beck, P. A., Intermediate Phases in

Binary Systems of Certain Transition Elements, Trans. AIME

206, p. 265 (1956).

2. Darby, J. B., Downey, J. W. and Norton, L. J., Intermediate Phases in the Tantalum- Palladium System, Trans.

AIME 227 , pp. 1028-1029 (1963).

3. Kane, R. , Giessen, B. C. and Grant, N. J., The Intermediate Phases AB, ABg and AB^ in the Ta-Pd and Ta-Pt Systems,

Journ. Metals (abstract), p. 687 Sept (1963).

4. Giessen, B. C. and Grant, N. J., Ne\v Intermediate Phases

. in Systems of Nb or Ta with Rh, Ir, Pd or Pt, Acta Cryst-17,

pp. 615-616 (1964).

5. Maldonado, A. and Schubert, K. Structure Investigation in

5 10

the T - T Homologous and Quasihomologous Alloy Systems,

Z. Met. 55, pp. 619-626 (1964).

6. Savitskii, E. M., Polyakova, V. P., Tylkina, M. A. and

Burkhanov, G. S. The Palladium- Tantalum System, Russian

Journal of Inorganic Chemistry 9^ pp. 890-893 (1964).

7. Waterstra^R. M. , The Vanadium -Platinum Constitution

Diagram, Met. Trans. 4, pp. 455-466 (1973).

8. Koster, W. and Haehl, W. D. , The Binary System PalladiumVanadium, Z. Met. 49, pp. 647-649 (1958).

9. Parker, D. ; Giessen, B. C. and Waterstrat, R. M. , The

Niobium- Palladium Phase Diagram, unpublished results

presented elsewhere in this report (1973).

10. Waterstrat, R. M., unpublished results (1973).

143
APPENDIX I: METALLOGRAPHIC PROCEDURES

Alloy specimens used for this study were in the form of

small pieces which had been removed from the arc- melted

buttons by an abrasive cut-off wheel or, in the case of brittle

alloys, by breaking the button with a hammer. Sufficient material was usually available to permit the use of separate pieces

for the various heat-treatments and consequently, remounting of

specimens was seldom necessary. Samples were, therefore,

mounted in bakelite. A compressed ball of copper turnings was

placed in the bakelite wrapped around the rear of the specimen

to produce a good electrical contact. A small hole was then

drilled into the mount from behind the specimen and into the

region containing the copper turnings but not deep enough to contact the sample directly. This hole served as an electrical

contact area for electro polishing or etching.

Grinding operations were performed manually on a

Handimet* surface grinder lubricated with cold running water.

Silicon carbide papers were used in the grit sequence 240, 320,

400 and 600. The samples were then given a preliminary polish

on Microcut paper sheets (Grit 600 soft). Final polishing procedures involved the use of 7 micron size diamond compound on a

Struers DUR cloth followed by 1/4 micron size diamond compound

on a Struers MOL cloth using the recommended lubricants. In

most cases, this was sufficient preparation for electro etching but in

some cases a final polish of MgO slurry on a Rayvel cloth was used.

* Certain commercial equipment, instruments, or materials

are identified in this paper in order to adequately specify the

experimental procedure. In no case does such identification

imply recommendation or endorsement by the National Bureau

of Standards, nor does it imply that the material or equipment

identified is necessarily the best available for the purpose.

144
The alloys used in this study are frequently very resistant

to conventional etchants such as hydrochloric or nitric acid and

consequently, it was necessary to develop more suitable methods

of etching. We have found that in most cases one may obtain

satisfactory micros true tures on these alloys by electro etching

in an aqueous solution of 2% to 10% potassium cyanide using 1 to

2 amperes at an applied AC voltage of up to 30 volts. This

method was particularly effective when used with a simultaneous

etch-polish technique which tended to reduce any tendency toward

pitting. The electrolytic cyanide etch is quite effective in delineating the boundaries of the various phases in these noble- metal

alloys and also in revealing grain boundaries or twinning. However, it frequently does not produce a contrasting color vvithin the

various phase regions and consequently, it is sometimes difficult

to establish the identity of an observed phase. In such cases the

use of the electron microprobe was most helpful since we were

unable to develop conventional etching procedures which would

produce color contrast in alloys having a high noble- metal content.

In a few cases, particularly in alloys having a high rhodium

or palladium content, it was not possible to obtain satisfactory

results using the cyanide etch. For these alloys we, therefore,

used a concentrated hydrochloric acid etchant electrolytically

with varying voltages and currents depending on the sample

compos ition.

Tantalum-palladium alloys were frequently etched with a

solution containing 1 part HF, 1 part HNOg, 2 parts H2SO^ and 5

parts H2O. This etchant was particularly effective in revealing

the microstructures in this system.

Niobium- rich alloys were etched with a 50% HF solution which

was particularly useful in the case of the niobium- palladium alloys.

145
APPENDIX II: X-RAY DIFFRACTION PROCEDURES

Debye-Scherrer powder patterns were taken of nearly all

equilibrated samples. A Phillips 114. 6 mm diameter camera

was employed for these pictures using nickel- filtered copper

radiation. In some cases, it was necessary to obtain an x-ray

pattern directly from the surface of a flat sample or to obtain

quantitative information on relative intensities. In such cases,

an x-ray diffractometer was used and Mr. Howard Swanson of

the National Bureau of Standards was most generous in assisting

us to obtain this data

Most of the lattice parameters given in this report were

obtained by carefully measuring the Debye-Scherrer films and

then applying the usual extrapolation methods in order to obtain

higher accuracy. However, the precision obtainable by the

methods which we employed is probably not higher than about

one part in 1, 000 but this seemed adequate for our purposes and

limitations in the available time prevented us from extend ii^

these measurements using methods of higher precision. In

most cases relative intensities were estimated by visual comparison of the lines on the Debye-Scherrer films.

Powder from the brittle alloys could be obtained simply by

crushing the alloy fragments in a glazed mullite mortar and

pestle but the more ductile alloys could not be crushed in this

manner. Powders of the ductile alloys were, therefore, obtained

by grinding a sectioned surface of the alloy fragment. Grinding

of these surfaces was accomplished using a dental diamond disc

(S. S. White #142 or 182) mounted in a horizontal 1/6 H. P.

electric motor rotating at 3500 rpm . During the grinding process the alloy powders became mixed with diamond particles

which had been dislodged from the discs. Fortunately the

diamond particles could be almost completely removed from the

collected powders by sieving through a 400- mesh screen.

146
The diamond particles remained on the screen while almost all

of the alloy powder passed through the screen. Therefore,

differential sieving, whereby the balance of the phase content is

altered, was apparently not a significant factor in our x-ray data.

If any residual diamond particles remained in the alloy powder,

it did not seriously interfere with our interpretation of the powder

patterns since the diamond lines were very spotty and readily

identified as such. Care was taken to replace the discs before

excessive wear had occurred since one would otherwise introduce

contamination from the substance used in bonding the diamond

powder to the metal disc.

Alloy powders obtained in this manner were highly strained

and contained much residual cold work. This would produce

considerable x-ray line broadening and such patterns would be

unsuitable for accurate lattice parameter measurements. It was,

therefore, necessary to remove the residual cold work by means

of a high- temperature annealing treatment. This was done by

placing the powder in a container of either Ta foil, alumina or

thoria, depending on the alloy composition, and then heating the

-5 powder in our vacuum furnace at a pressure of less than 10 mm

for a few seconds to several minutes depending on alloy composition and annealing temperature. In all cases the annealing

temperature used for the powders was identical to that used in the

last equilibration treatment for the bulk sample from which the

powder had been obtained. After the annealing treatment, the

powders were rapidly cooled by turning off the furnace power.

Excessive annealing resulted in a partial sintering of the powders

and, therefore, the minimum annealing time needed to effectively

remove the internal strains is usually employed. Satisfactory

x-ray diffraction patterns were obtained using these procedures.

147
APPENDIX m: ELECTRON MICROPROBE ANALYSIS

The electron microprobe is an instrument capable of

focusing a beam of electrons on a spot approximately one micron

(0. 00004 inch) in diameter. These electrons excite x-rays,

characteristic of the elements irradiated, which can be analyzed

to yield a quantitative chemical analysis of the spot irradiated.

The probe is equipped with an optical microscope and the electron

beam is focused in the desired areas. The wave lengths emitted

by the elements in the specimens are separated by diffraction from

an analyzing crystal, and their intensities are recorded by radiation

counters. By comparing the observed intensity of characteristic

radiation from the sample with that of a pure element as a standard,

a quantitative analysis can be obtained for the irradiated spots.

The characteristic intensities for a given element in the

sample and in the standard are first corrected for pulse coincidence losses and background. From this data one obtains an

intensity ratio for the sample relative to the standard. The

percentage of a given element necessary to produce such an

intensity ratio is calculated by either of the programs COR or

MULTI8 on a digital computer. This introduces corrections for

x-ray absorption, atomic number differences and secondary x-ray

fluorescence. The calculations are checked by analyzing alloys of

known composition which have been chemically analyzed.

Phase identity was readily established for each sample by

operating the probe as a scanning electron microscope. This

permits one to establish which of the two phases has the highest

average atomic number as revealed by differences in brightness

of the specimen image on an oscilloscope screen. This was

particularly helpful in our alloys since the identity of each phase

148
could not usually be clearly established by chemical etching

methods but differences in the atomic numbers were quite large.

The determination of phase diagrams using annealed two-phase

alloys is particularly efficient when each individual phase can be

chemically analyzed. Since the phases are in equilibrium with

one another, the phase boundary compositions are given directly

by the compositions of the two phases. This permits a great

saving in the number of samples needed. Only one sample is

required for each equilibration annealing temperature in a given

two-phase region. Conventional methods of metallographic examination usually require the preparation of many different alloy

compositions in order to obtain an identical amount of information

by bracketing the phase boundary. The use of the electron microprobe also provides valuable information on the equilibrium status

of the alloys and the effectiveness of quenches, since it very quickly

reveals composition gradients or precipitation of a third phase.

We were particularly careful in searching for fine

precipitated particles in our samples since a representative microprobe analysis will not be obtained when such precipitates occupy

an area smaller than the diameter of the electron beam. In several

cases, fine precipitates were suggested as being responsible for

erratic results during microprobe analysis. The presence of these

precipitates was later verified by optical microscopy. It is not

possible to obtain an accurate quantitative analysis from phase

regions containing such fine precipitates.

149
APPENDIX IV: ATOMIC VOLUME RELATIONSfflPS

It is sometimes desirable to compare the unit cell

dimensions for each of the various phases which occur in a given

alloy system. In order to make such a comparison it is necessary

to bring the unit cell dimensions to a common basis. This is

done by plotting the lattice parameter data in the form of atomic

volumes (volume of the unit cell divided by the number of atoms

per unit cell) as a function of composition. If one draws a

straight line between the atomic volumes of each elemental

component, then it is possible to determine whether any net

expansion or contraction has occurred in the various intermediate

phases, i.e., whether there is a positive or negative departure

from a "volume Vegard's Law" relationship.

Plots of atomic volume versus composition have, therefore,

been constructed for the phases occurring in each of the six

binary systems explored in this investigation. These plots are

shown in Figs. 37-42.

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