NATIONAL BUREAU OF STANDARDS REPORT

10 571

Progress Report

on NOBLE METAL CONSTITUTION DIAGRAMS

U.S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS
NATIONAL BUREAU OF STANDARDS

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NATIONAL BUREAU OF STANDARDS REPORT

MBS PROJECT MBS REPORT

311.05—11—3110561 April 20, 1971 10 571

Progress Report

on NOBLE METAL CONSTITUTION DIAGRAMS

by

R. M. Waterstrat and R. C. Manuszewski*

* Research Associates, American Dental Association

Research Program in the Dental Research Section,

National Bureau of Standards, Washington, D. C. 20234

This investigation was supported in part by Research Grant

DE02455 to the American Dental Association from 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

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

and Development Command; the Dental Sciences Division of the School of Aerospace Medicine, USAF; the National Institute

of Dental Research; and the Veterans Administration.

IMPORTANT NOTICE NATIONAL BUREAU OF ST I for use within the Government, and review. For this reason, the whole or in part, is not authori Bureau of Standards, Washingto

the Report has been specifically

Approved for public release by the

;s accounting documents intended

subjected to additional evaluation

listing of this Report, either in Director of the National Institute of

Standards and Technology (NIST) on October 9, 2015.

: Office of the Director, National

y the Government agency for which

opies for its own use.

U.S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS

October 15, 1970

NOBLE METAL CONSTITUTION DIAGRAMS

by

R. M. Waterstrat and R. C. Manuszewski*

Research Associates, American Dental Association

Research Program in the Dental Research Section,

National Bureau of Standards, Washington, D. C 0 20234

This investigation was supported in part by Research Grant DEQ2455

to the American Dental Association from 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 Dental Research Division of the U. S. Army Medical

Research and Development Command; the Dental Sciences Division of

the School of Aerospace Medicine, USAF; the National Institute of

Dental Research; and the Veterans Administration.

i
NBS PROGRESS REPORT 10 571

NOBLE METAL CONSTITUTION DIAGRAMS

R. M. Waterstrat and R„ C. Manuszewski

SEPTEMBER 1970

Performed under

NIDR Research Grant DE02455 to

The American Dental Association.

Conducted at the

National Bureau of Standards

Washington, D. C. 20234

li
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, 1966 to August 31, 1970.

The electron probe microanalyses of two-phase alloys was

performed at the National Bureau of Standards by Mr. Don Vieth,

Mr. Barry Hammond, Mr. Robert Myklebust and Mr. Charles Fiore

under the direction of Dr. Kurt Heinrich.

iii
ABSTRACT

Four binary constitution diagrams involving the noble

metals are presented. These diagrams include the Cr-Pt,

Cr-Ir, Cr-Rh and V-Pt 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

A. Constitution Diagrams . . 1 B. Experimental Method . . 2

H. INTRODUCTION 4

A. Background 4

IE. CHROMIUM-PLATINUM CONSTITUTION DIAGRAM 9

A. Previous Studies 9

B. Materials 10

C. Alloy Preparation .... .... 12

D . Composition Determination of the Melts 15

E. Temperature Measurement . 16

F. Thermal Treatments 17

G. Methods for Determination of Phase Boundaries 18

H. Experimental Results 26

REFERENCES 44

IV. VANADIUM -PLATINUM CONSTITUTION DIAGRAM ..... 46

A. Previous Studies . 46

B. Materials 46

C. Alloy Preparation 47

D. Composition Determination of the Melts 48

E. Temperature Measurement 49

F. Thermal Treatments 49

G. Methods for Determination of Phase Boundaries 51

H. Experimental Results 54

REFERENCES. 98

v
TABLE OF CONTENTS (Cont'd.

)

Page No.

V. CHROMIUM-IRIDIUM CONSTITUTION DIAGRAM 100

A. Previous Studies 100

B. Materials 100

C. Alloy Preparation 101

D. Composition Determination of the Melts 103

E. Temperature Measurements 103

F. Thermal Treatments 104

G. Methods for Determination of Phase Boundaries .... 106

H. Experimental Results Ill

REFERENCES 124

VI. CHROMIUM-RHODIUM CONSTITUTION DIAGRAM ..... 125

A. Previous Studies 125

B. Materials 125

C. Alloy Preparation . 126

D. Composition Determination of the Melts 127

E. Temperature Measurements 128

F. Thermal Treatments 129

G. Methods for Determination of Phase Boundaries .... 130

H. Experimental Results 134

REFERENCES . 151

vi
LIST OF ILLUSTRATIONS

Figure Page No.

1. Constitution diagrams of Cr-Pt, V-Pt, Cr-Ir and Cr-Rh . . 1 2. Apparatus used in the determination of solidus temperatures . 19

3. Vacuum furnace hot-zone temperature as a function of

heater current „ 24

4. The chromium-platinum constitution diagram 27

5. Microstructures of chromium-platinum alloys 33

6. Variation in the lattice parameters of Cr-Pt alloys as a

function of alloy composition 34

7. Microstructures of chromium-platinum alloys 35

8. Microstructures of chromium -platinum alloys 36

9. Microstructures of chromium -platinum alloys 37

10. Microstructures of chromium -platinum alloys 38

11. Microstructures of chromium-platinum alloys 39

12. Microstructures of chromium-platinum alloys 40

13. The vanadium -platinum constitution diagram 55

14. Relationship of the Al3 Ti-type and Cus Au-type structures. . 64

15. Microstructures of vanadium -platinum alloys 67

16. Microstructures of vanadium -platinum alloys 72

17. Thermal expansion of VPt as a function of temperature . . 75

18. Microstructures of vanadium -platinum alloys ..... 77

19. Microstructures of vanadium -platinum alloys 78

20. Microstructures of vanadium -platinum alloys 79

21. Powder x-ray diffraction patterns of single phase V-Pt

alloys obtained with nickel-filtered copper radiation ... 88

22. Microstructures of vanadium -platinum alloys ..... 90

23. Microstructures of vanadium -platinum alloys 91

24. Microstructures of vanadium -platinum alloys 92

25. Microstructures of vanadium -platinum alloys 94

26. Microstructures of vanadium -platinum alloys ..... 95

27. Microstructures of vanadium -platinum alloys ..... 96

28. Microstructures of vanadium -platinum alloys 97

29. The chromium-iridium constitution diagram ..... 112

vii
LIST OF ILLUSTRATIONS (Cont’d.)

Figure Page No.

30. Microstructures of chromium -iridium alloys 113

31. Lattice parameters of € phase Cr-Ir alloys 116

32. Microstructures of chromium-iridium alloys 118

33. Microstructures of chromium-iridium alloys 119

34. Microstructures of chromium -iridium alloys 120

35. Microstructures of chromium-iridium alloys 121

36. Microstructures of chromium -iridium alloys 122

37. The chromium-rhodium constitution diagram 135

38. Lattice parameters of € phase Cr-Rh alloys 139

39. Microstructures of chromium-rhodium alloys 140

40. Microstructures of chromium -rhodium alloys 141

41. Microstructures of chromium -rhodium alloys 142

42. Microstructures of chromium -rhodium alloys 143

43. Microstructures of chromium-rhodium alloys 144

44. Microstructures of chromium -rhodium alloys . .... 145

45. Microstructures of chromium -rhodium alloys 146

46. Microstructures of chromium-rhodium alloys 147

47. Microstructures of chromium-rhodium alloys 148

48. Schematic diagram of high-temperature vacuum furnace . 159

49. High -temperature vacuum furnace in operation .... 161

50. Atomic volumes of Cr-Pt alloys 163

51. Atomic volumes of V-Pt alloys ........ . 164

52. Atomic volumes of Cr-Ir alloys 165

53. Atomic volumes of Cr-Rh alloys . 166

viii
LIST OF TABLES

Page No.

1. Chemical and Spectrographie Analyses of the Materials Used

in Preparing the Experimental (Cr-Pt) Alloys 11

2. Analyses of Inert Gases Used for Arc -melting the

Experimental Alloys 14

3. Chemical Analyses of Experimental (Cr-Pt) Alloys . 15

4. Summary of Equilibration Treatments for ChromiumPlatinum Alloys . 29

5. Typical Analyses of the Materials Used in Preparing

the Experimental (V-Pt) Alloys . . . . . 47

6c Summary of Equilibration Treatments for Vanadium

Platinum Alloys ....... ....... . 58

7. Lattice Parameters of y Phase V -Pt Alloys ...... 61

8. X-ray Diffraction Pattern of 0 Phase VPt3 ....... 65

9o X-ray Diffraction Pattern of 6 Phase VPt2 ....... 69

10. X-ray Diffraction Pattern of C Phase VPt . 81

1L X-ray Diffraction Pattern of p Phase V3 Pt ....... 86

12. Chemical and Spectrographie Analyses of the Materials Used

in Preparing the Experimental (Cr-Ir) Alloys ...... 102

13. Summary of Equilibration Treatments for Chromium

Iridium Alloys 107

14. Comparison of Cr-Ir Alloy Compositions Determined by

Weight Loss and by Electron Probe Microanalysis .... 110

15. Chemical and Spectrographie Analyses of the Materials Used

in Preparing the Experimental Alloys 127

16. Summary of Equilibration Treatments for Chromium

Rhodium Alloys . .............. 131

ix

I SUMMARY

A. Constitution Diagrams

The constitution

are presented in Figure 1.

diagrams which have been determined

WEIGHT PERCENT CHROMIUM

Fig. 1 Constitution diagrams of chromium- platinum; vanadium -platinum

chromium-iridium; and chromium-rhodium.

1
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. Chromium was in the form

of fused droplets of about 1/8-inch diameter. Vanadium was in the

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

diameter cylinders of the appropriate weight. Iridium and rhodium

powder or sponge was also compressed into 1/2 -inch diameter

cylinders. Platinum was in the form of 1/32 -inch sheet 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 150Q°C and 220Q°C were measured using tungsten versus

2
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 iridium0 The

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

5. Phase Boundary Determinations

a. Solidus - A technique was developed to

allow direct observation of a melting specimen and 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

metallo graphic 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.

3
II. INTRODUCTION

A0 Background

The existing metallurgical literature concerning binary

or ternary constitution diagrams contains only a relatively small percentage of alloy combinations involving the noble metals (i.e., those

metals which are exceptionally resistant to chemical corrosion or

oxidation, such as gold and metals of the platinum group). 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 microstructures 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

non- 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). A knowledge of the solidus temperature permits one to estimate recovery

temperatures, recrystallization temperatures, and the rates of atomic

4
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.

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 high

accuracy and reproducibility associated with this type of thermocouple. These thermocouples were periodically checked for

accuracy against a standard platinum-6% rhodium versus piatinum30% rhodium thermocouple which had been calibrated at the

National Bureau of Standards to 1600° C.

5
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

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 tungsfen-26% rhenium thermocouple insulated with

beryllium oxide and having a tantalum sheath0 The size of the

thermocouple wires was 0.005" diameter and the entire assembly

was enclosed in a sheath having a diameter of 0.040", thus minimizing errors due to heat conduction along the axis of the thermocouple.

6
The tungsten versus tungsten-26% rhenium thermocouples

were calibrated against the standard platinum- 6% rhodium versus

platinum- 30% rhodium thermocouple to 1600°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 deter mination0 This procedure is explained in

greater detail in Section III, Q-l.

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

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 usually

included 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. Metaliographic studies, particularly

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

detecting contamination.

7
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.

8
III. CHROMIUM- PLATINUM CONSTITUTION DIAGRAM

A. Previous Studies

Chromium-platinum alloys have been studied in

several previous investigations but the results have been

(1-5) contradictory. In reviewing these studies, v ’ Hansen and

Anderko

^

have concluded that the data ’’are still insufficient

to establish the exact constitution oj: the system.”

' ( 1 ) Melting point data obtained by Muller ' ' have indicated that the addition of chromium to pure platinum produces

a broad maximum in the liquidus but thermal analysis data from

three chromium-rich alloys

^

has led Gebhardt and Koster to

conclude that Muller’s melting curve is much too high, Hansen

and Anderko

^

have also concluded that Muller’s melting points

’’are certainly too high.”

The present study contradicts these conclusions and

establishes a location of the solidus which is in substantial agreement with Muller’s results. The solidus temperatures are found

to be significantly higher than those reported by Gebhardt and

(4) Koster. v ' Other results of Gebhardt and Koster have been con-

(5) tradicted by Raub and Mahler v ' who reported that the platinum

solid solution and the chromium solid solution are separated by

an intermediate phase, PtCr Q , having the A15 ( $-W) crystal

0

(4) structure. Gebhardt and Koster v ' apparently did not detect

this intermediate phase and consequently they erroneously concluded that the two terminal solid solutions coexist in the region

between 80 and 100 atomic percent Cr. The present results are

(5) in agreement with those reported by Raub and Mahler v ’ except

that the intermediate phase is found to occur only for compositions

on the Cr-rich side of the "ideal” (PtCr^) stoichiometry.

9
B. Materials

The chromium metal used in this investigation was

obtained from Leico Industries, Incorporated in the form of

small droplets of about 1/ 8-inch diameter which were quite

suitable for arc- melting. It’s nominal purity was 99.999%

with respect to nongaseous impurities. A spectrographic

analysis of this material is shown in Table 1. The gaseous

impurities were determined by vacuum fusion analyses and

these results are also given in Table 1. It is apparent that the

major impurity in this material is oxygen, which is present at

a rather high concentration of over two atomic percent. This

is an appreciable impurity content, and one must consider

whether it can be responsible for significant changes in the

location of phase boundaries. This problem was dealt with

in two different ways.

One approach was to melt the alloys using the chromium

with high oxygen content and then to deoxidize these alloys by

annealing them in dry hydrogen at temperatures of 1400° or

above. The effect of such a deoxidation anneal is quite visible

in the microstructures of these alloys. It is manifested by a removal of oxide particles (presumably Cr203 ) from the microstructure. The oxide particles apparently contain most of the

oxygen in these alloys. It has been reported that oxygen is

not very soluble in chromium at low temperatures. One may

therefore conclude that those samples which were relatively

free of oxide inclusions were also essentially free of oxygen

and that they can be used to establish the location of phase

boundaries in the binary system. A comparison of experimental

data obtained from the hydrogen- treated alloys with data obtained

from the oxygen- containing alloys revealed that the location of

phase boundaries was essentially unchanged by the observed

oxygen contamination.

10
Table 1

Chemical and Spectrographic Analyses of the Materials

Used in Preparing the Experimental Alloys +

Chromium Platinum

Impurity Leico Metals Inc. Stock supply

Ag ND <10

Ca <10 ND

Cu ND <100

Fe <10 <10

Mg <10 <100

Ni <10 ND

Pb <10 ND

Pd ND <1000

Ru ND <100

Si <10 <100 H*C 16

* 295

O* 6540

N* 6

t Values are given in ppme ND Not detected spectrographically0

* Values obtained by vacuum fusion analysis,.

No other elements were detected spectrographically.

11
Another approach to the problem of oxygen contamination was to submit the pure chromium pellets to a heat- treatment

in dry hydrogen prior to melting the experimental alloys. These

heat- treatments were performed over various time intervals

ranging from three hours at 1700° C to three days at 1400° C.

Alloys prepared from the deoxidized chromium pellets yielded

results similar to those obtained from samples prepared from

untreated chromium. For example, the microstructure of our

untreated alloy, Cr 80Pt20» in the ’’as-cast" condition (Fig. Ill}

has been largely responsible for our conclusion that a syntectic

reaction occurs in the chromium-platinum system. This same

microstructure has been obtained in "as-cast" alloys prepared

using deoxidized chromium.

We have therefore concluded that oxygen contamination

to the extent of up to two atomic percent does not change the

location of the phase boundaries within the experimental error

of our measurements.

The platinum metal was available from a stock supply

whose original source was unknown. Spectrographic analysis,

shown in Table 1, indicates a purity of about 99. 9% with the

major impurity being palladium. The platinum metal was in

the form of 1/32- inch thick sheet which could be readily cut

and bent to a shape which is suitable for arc- melting.

C. Alloy Preparation

The experimental alloys were prepared by weighing

the appropriate amount of each elemental constituent and then

melting these constituents in an arc furnace under an inert

atmosphere. Further details are given below.

12
1. Weighing ,

Samples were weighed out in 10 atom percent

intervals across the entire phase diagram except in the vicinity

of the anticipated phase boundaries where additional compositions

were prepared. The components of each alloy were weighed to an

accuracy of ± 0. 01 gram and the total weight of each sample was

approximately 20 grams. It was therefore possible to establish

the nominal compositions of our alloys to within ± 0.05 weight

percent. The weights were checked to this accuracy at various

stages in the sample preparation in order to obtain a material

balance for each alloy which would permit the establishment of

maximum uncertainties in the alloy compositions,,

2. Melting

The samples were arc melted on a water-cooled

copper hearth using a non-consumable tungsten electrode 0 All

melting was done in an atmosphere of 50% helium and 50% argonQ These inert gases were of a special ultra-pure grade for which the

gas analyses shown in Table 2 were provided by the manufacturer.

No further purification of these gases was attempted but careful

precautions were taken to insure that gas lines and regulators

were clean and leak-tight in order to prevent contamination. The

single- hearth melting chamber was evacuated and flushed several

times with the inert gas mixture and the final evacuation was

3 extended to a pressure of 10“ mm or less. This pressure corresponds to a residual impurity level of less than 3. 0 ppm which is

comparable to the residual impurity levels in our ultra-pure inert

gases as shown in Table 2 0

Each sample was melted at least four times and

was inverted in the hearth between each melt. Special precautions

were taken with brittle alloys to insure that no fragments of the

alloy were lost if a sample shattered under the arc. It was found

that brittle alloys were much less susceptible to such losses if the

13
Table 2

Analyses of Inert Gases Used for

Arc- melting the Experimental Alloys +

Argon Helium

Impurity Matheson Co. Matheson Co.

co2

ND* 0.5 0.0

°2 1.0 0.3

h2

NB* 1.0 0.0

CO ND* 0o 5 -- N

2

< 2.0 1.0

ch

4

< 0.4 0.0

h

2

o 3. 5 1.5

Ne — 9.2

A --- 0.0

t Impurities are given in ppm.

* None detected less than.

alloys were broken up into irregular lumps of about 1/4- inch thickness prior to any subsequent re melting. This was done using a hardened steel rod in a deep cylindrical hardened steel container

so that none of the alloy fragments were lost. For most of the

brittle alloys, only a light sharp blow was necessary to fracture

the sample and no contamination from the steel container was

detected.

14
D. Composition Determination of the Melts

The alloy compositions were established by chemical

analyses, electron probe microanalyses, weight balances and

lattice parameter measurements. Weight losses which occurred

during melting of the platinum-rich alloys were always less than

1% but some chromium-rich alloys exhibited weight losses as high

as 1.2% after four melts, or 1.7% after six melts. It was subsequently found that good agreement with the chemical analyses of

our alloys was obtained if these weight losses were attributed

entirely to losses of chromium0

Table 3

Chemical Analyses of Experimental Alloys

Intended

Composition

Chromium

atomic %

Platinum

atomic %

Cr

10

Pt90 9.7 90.2

Cr2Q

Pt

80 20. 0 80.0

Cr

3G

Pt

70

29 0 7 70.4

Cr

4Q

Pt

6Q 39. 5 60.4

Cr

5Q

Pt

50 49.7 50.2

Cr60

Pt

40 59.4 40.5

The chemical analyses, shown in Table 3, were

obtained from alloy samples which had been given a prior homogenization anneal in vacuum at elevated temperatures. The annealing treatment produced small weight losses and a metallographic

15
study revealed that these losses were essentially confined to a

very thin layer on the surface of the sample which haa been

depleted of chromium. The homogenisation anneal was given

prior to chemical analysis in order that the chemical analysis

results could be correlated with the electron probe data Th

electron probe microanalyzer measures areas of about o

micron diameter so the data would be unreliable if segrega ion

occurred on a larger scale than this. However, it wassu

quently found by line scanning that our platinum-rich alloys

were quite homogeneous, even in the "as-cast” condition. This

may be due to the existence of a congruent melting maximum

the platinum solid solution region.

Each of the alloys shown in Table 3 was cut in half

after the homogenization anneal. One half was submitted 10

Chemical analysis and the other half was used to obtain ^ec on

microprobe data. The results of the chemical analyses thus

established the compositions of the standard samples which were

used in calibrating the electron probe data on a quantitative basis.

The lattice parameters of alloys within the range of

the platinum solid solution are plotted as a function of composition

in Fig 6 and the chemically analyzed alloys are included in

this plot. The variation of the lattice parameters with composition

is linear and the experimental data is in excellent agreemen

data reported by Raub and Mahler ^ > which is also plotted in

Fig. 6

E Tem.pera.ture Measurement

A platinum versus platinum-10% rhodium thermocouple

was used to measure temperatures up to 1500°C. Temperatures

above 1500°C were measured with a tungsten versus -ungs en o

rhenium thermocouple. Both of these thermocouples were peno -

ically calibrated against a standard platinum-o % rhodium vers

16
platinum-30% rhodium thermocouple at temperatures up to 1600° C.

The standard thermocouple had been calibrated at the National

Bureau of Standards. In addition, the thermocouples were calibrated against secondary standards such as the melting points of pure

nickel or pure platinum, particularly in connection with the determination of solidus or liquidus temperatures. Details of this

procedure are given in Section HI, G.

F. Thermal Treatments

1. Homogenization

Most of the samples used in this study were

initially homogenized by heating to within about 200° C of the

solidus temperature for about three hours0 The homogenization

treatment was conducted with all the precautions normally used

in the equilibration treatments so that the samples could be considered as being equilibrated at the homogenization temperature.

Samples to be homogenized were placed in

tantalum buckets suspended by tantalum wire within the hot zone

of a tantalum strip resistance furnace operating at a pressure of

between 5 x 10“° and 5 x 10”^ mm Hg. No reaction between the

samples and the tantalum was detected, and it is probable that

any small reaction with Ta is insignificant in the temperature

ranges which were used in this study.

The degree of homogenization was checked by

means of the electron probe microanalyzer, metallographic

examination, and by the sharpness of x-ray powder patterns

obtained from the homogenized samples. In all cases, the

samples appeared to be adequately homogenized,,

17
2. Equilibration

Equilibration anneals at temperatures up to

1100° C were conducted by sealing the samples in quartz tubes

(previously baked-out and evacuated) under a partial pressure

of pure argon and then placing these tubes in a platinum- wound

resistance furnace capable of being controlled within a temperature range of ± 1°C. For temperatures above 1100°C, the

samples were equilibrated in a tantalum- strip resistance furnace operating at a pressure of 5 x 10“ ® to 5 x 10“^ mm Hg.

The samples were suspended in tantalum buckets but no tantalum

pickup was detected in the alloy samples 0 Temperatures in

this furnace were normally controlled to within ± 10° C, but

could be maintained within ± 5°C over relatively short periods

of time. This was verified by measuring temperature profiles

through the hot zone using thermocouples which were inserted

both vertically and horizontally. The temperature profiles

were also used to establish the location of various temperature

gradients within the hot zone. The alloy samples could then be

located within the constant- temperature region.

Samples annealed in the tantalum- strip furnace

were cooled rapidly from the annealing temperature by turning

off the power to the heating elements. Samples annealed in the

quartz tubes were quenched in cold water. In each case, the

cooling rates were sufficient to permit retention of the phases

which existed at the annealing temperature. However, atomic

ordering in the Pt terminal solid solution region occurs very

rapidly and it was not possible to retain the disordered condition

for alloys containing over 20 atomic percent Cr.

G. Methods for Dete rmination of Phase Boundaries

1. Solidus Temperatures

The accurate measurement of solidus temperature by conventional thermal analysis was found to be extremely

18
TUNGSTEN

vs.

TUNGSTEN-26%

RHENIUM

THERMOCOUPLE

WIRES

—

O

o

19

Fig'.

2

-

Apparatus

used

in

the

determination

of

solidus

temperatures.
difficult due to the high vapor pressures and chemical activity

of chromium when it is heated to relatively high temperatures.

Thermocouples tend to become contaminated by Cr vapor, which

may even penetrate the refractory tubes used as insulators for

the thermocouple. Many refractories are chemically attacked

by molten chromium alloys with a consequent serious contamination of the alloy. Optical pyrometry techniques are unreliable

due to a rapid condensation of Cr vapor on the glass windows

through which the optical measurements are made. The use of

window shields, which are removable during the time interval

necessary for a temperature measurement, helps to reduce

fogging of the windows but does not appreciably alleviate the

problems of temperature measurement due to the rapid rate

at which the .vapors are deposited. We were, therefore, unable

to use conventional methods of thermal analysis in most cases,

and were forced to utilize a method which is more suited to

these severe experimental conditions.

A schematic drawing illustrating the technique

which was used is shown in Fig. 2. A previously homogenized

alloy specimen, together with a pure metal of known melting

temperature, were placed in suitable refractory crucibles (alumina

or thoria) in a manner such as to expose the upper portion of each

metal to direct visual observation along the direction indicated by

the arrows in Fig. 2 . The crucibles are inserted into cylindrical holes drilled part- way into a tantalum cylinder of i-l/4-inch

diameter and about 7/8-inch height. The tantalum cylinder is

suspended vertically within the constant temperature region of

the vacuum furnace by means of tantalum wire which is inserted

into three holes drilled through the tantalum cylinder near its

periphery. At the center of the cylinder a vertical hole drilled

part-way through the cylinder admits a sheathed thermocouple of

tungsten versus tungsten- 26% rhenium which can be used to measure

temperatures at least up to about 2300° C in vacuum. The tantalum

thermocouples sheath has a diameter of about 0. 040M

, and is

20
closed at its lower end thus acting to protect the thermocouple

from attack by metal vapors. The beryllia insulator serves to

provide electrical insulation at high temperatures and is chemically inert to the thermocouple.

The alloy sample, which preferably has a jagged

edge, and the standard wire are visible against a relatively cool

background as viewed between the tantalum strips comprising the

heating element of the furnace. Melting is readily detected by

visual observation using dark glasses at the higher temperatures.

The upper portions of the standard wires were found to melt before

the lower portions melted. It was necessary to increase the furnace temperature by about 10° C in order to melt the lower portions

of the wires. The alloy samples, however, extend only a short

distance into the crucibles and when an alloy sample has melted,

one generally observes that the entire sample suddenly falls down

into the crucible and solidifies 0 The tip of the thermocouple is

located within the tantalum cylinder near the base of the crucibles,

however, and will probably measure a temperature approximately

10° C lower than that of the alloy samples which are visible above

the rim of the crucibles 0 This correction, however, is included

in our temperature calibration procedure since a thermocouple

reading is obtained as the top of the standard wire begins to melt.

The top of the wire is probably at a temperature nearly equal to

that of the alloy specimen. Measurement of a vertical temperature profile can also be made by raising and lowering the thermocouple. This procedure also helps in establishing the temperature

difference between the sample and the normal location of the

thermocouple 0

Losses of metal by vaporization during a solidus

determination occurred mainly at the surface of the alloy specimens.

Lattice parameter measurements on the bulk sample after a run

usually indicated a composition change of about 1% or less 0 It appears that appreciable losses due to vaporization occur only after

21
the formation of a liquid phase. Our samples are usually in the liquid

condition for only a few seconds during the run, thus minimizing

metal losses. The use of a high vacuum during solidus determinations, therefore, seems justified since the avoidance of possible

sample contamination by gaseous impurities prevents more serious

errors than are introduced by the metal losses.

When making a melting temperature determination, the temperature of the furnace was raised in increments of

about 5°C as the solidus point was approached. Between each

increment, the temperature was allowed to equilibrate by waiting

for the temperature -sensitive thermocouple to indicate a constant

temperature.

This usually required less than a minute near

the solidus temperature. Meanwhile, the metal samples are

checked visually for signs of melting. The pure metals and many

of the alloy samples melted very suddenly, but some of the alloy

samples showed only slight rounding of sharp edges or changes in

brightness as the solidus temperature was exceeded. Any question

as to whether the sample had actually begun to melt could usually

be settled by metallographic examination of the sample after its

removal from the furnace. The thermocouple wires were of 0o QQ5"

diameter and it is, therefore, assumed that the temperature reading is not greatly affected by heat losses along the thermocouple

wires 0 A possible source of error in using pure metals

with known melting points as temperature standards may arise if

the pure metals become contaminated during the calibration procedure. We found no evidence for any significant contamination

of the pure metal wires. Wires melted with and without the

chromium alloy specimens melted at nearly the same temperature as indicated by the thermocouple readings and values

of furnace current at tne melting point.

22
Our standard platinum-6% rhodium versus

platinum- 30% rhodium thermocouple was inserted in the furnace

under standard operating conditions and a calibration of temperature versus furnace current was obtained up to 1700° C. This

calibration was extended to higher temperatures by obtaining furnace current readings at the melting points of pure platinum

(1772°C), pure rhodium (1963°C), and pure iridium (2447° C). The

pure metal wires were of special high purity (99.99%) and had

unusually low Fe and Si contents (<100 ppm). The use of furnace

current as an indication of the temperature in the hot zone was

considered fairly reliable c The method yielded good reproducibility under standard operating conditions and reasonable accuracy

resulted owing to the fact that furnace temperature is a fairly

sensitive function of furnace current as shown in Fig. 3.

The technique of visual observation of melting

was considered to be best suited for pure elements and alloys

melting isothermally. Thus, it is most reliable in the vicinity

of congruent melting maxima and near eutectic compositions.

For alloys melting over a wide temperature range, however,

this method may yield erroneous results since such alloys may

show no visual evidence of melting, even though much melting

may have occurred. In such cases, the technique of ” incipient

melting” was used to obtain more accurate solidus temperatures.

In this technique, an alloy sample is heated for a few minutes at a selected temperature. The furnace power is then turned off

and the sample is removed and examined metallographically.

If the solidus temperature has been exceeded, one usually observes rounded primary grains with a network of fine grains

at the prior grain boundaries of the larger primary grains.

Comparison of this structure with structures obtained by heating

at lower temperatures enables one to determine the solidus

temperature to the desired accuracy within the limits of the

temperature calibration.

23
3000 o o o © ©

o o © o ©

ID o © in

OJ

Co •) lunivuimu

o

CO

yj

yj

a

2

<

sr

Of

o

Of

yj

<

yj

24

Fig.

3

-

Vacuum

furnace

hot-zone

temperature

asa

function

of

heater

current.
2 0 Liquidus Temperatures

Due to the previously mentioned experimental

difficulties of conducting thermal analyses on chromium alloys

at high temperatures, we have not obtained direct measurements

of liquidus temperatures. The approximate liquidus temperatures

were estimated from the metallographic appearance of samples

quenched from above the solidus. Given the form of the solidus

and the temperature at which the specimen was held, one can

estimate the liquidus composition by applying the lever rule after

estimating the relative amount of solid and liquid phases as

revealed by metallographic examination.

3. Invariant Reactions

The phases existing at various temperatures

were determined by metallographic examination and DebyeScherrer x-ray examination of alloys which had been annealed

at these temperatures followed by rapid cooling. From such

studies, one may characterize the structure existing above a

given transformation temperature and then compare this structure

with those obtained for alloys which have been annealed below the

transformation temperature. This technique can be used to

bracket the transformation temperature by choosing successive

annealing temperatures above and below the transformation

temperature. Such a procedure converges fairly rapidly and

one can usually determine the transformation temperature to

within the accuracy of the temperature measurement itself. By

using this technique, one may avoid the supercooling effects which

are a complicating factor when conventional thermal analysis

methods are employed.

Eutectic transformation temperatures may also

be determined using the methods described for the determination

of the solidus temperatures (Section III, G-l). After the solidus

25
determination the experimental sample is examined metallographicallye One can then estimate the eutectic composition by observing the amount of any typical eutectic structure relative to the

amount of coexisting phases of known compositions 0 The observation of a typical eutectic structure also helps to verify that the

transformation is indeed a eutectic transformation.

4. Other Phase Boundaries

The soivus lines and the phase boundaries of

intermediate phases were obtained by electron probe microanalysis

whenever it was possible to use this method. Details of the experimental procedures are given in the Appendix. In some instances, it

was not possible to use this method due to the presence of fine particles of precipitated phases or other finely decomposed structures.

In such cases, the diameter of the electron beam is too large to

permit its impingement entirely within a single phase region and

metallo graphic examination is, therefore, used to determine the

location of the phase boundary.

H. Experimental Results

1. General

The proposed constitution diagram for the

chromium-platinum system is presented in Fig. 4„ The various

features of this diagram are discussed below. All alloys were

inspected either metallographically or by means of x-ray, DebyeScherrer powder patterns, or both. In general, good agreement

exists among the metallographic, Debye -Schemer, and electron

microbeam probe data.

2. Solidus

The solidus was determined mainly by using the

technique of direct observation of melting using the cylindrical

26
WEIGHT

PERCENT

CHROMIUM

Jo ‘3dfl±Vd3d IAI31

O

crX

o

X

UJo

ccUJ

CLUJo

h-<

V

Oo ‘ 3dfl±Vd 3d IAJ31

27

Fig.

4

-

Chromium

-platinum

constitution

diagram
tantalum holder as shown in Fig 0 2, and described in Section HE,

G-l. In most cases, melting occurred quite rapidly at a definite

temperature level, particularly in the v solid solution region

where a congruent melting maximum presumably exists. Each

solidus determination included an alloy sample and also a piece of

high purity platinum wire so that a simultaneous temperature

calibration was obtained at the Pt-point.

This technique seems quite reliable since the

Pt-point calibration is done during the course of the actual solidus

determination and under identical experimental conditions,, The

sheathed thermocouple is quite sensitive to temperature changes

of the order of 1°C or less, and it was, therefore, quite easy to

detect the small temperature difference which usually existed

between the observed melting point of the platinum wires and

those of the alloy samples. High purity ARH alumina crucibles

were used in these experiments. The crucibles appeared to

resist any chemical attack except in the case of Cr-rich alloys

where a slight green discoloration was detected. These samples

were rerun using thoria crucibles, but we observed no significant differences in our results between the use of alumina versus

thoria crucibles for the Cr-rich alloys.

3. Equilibration

Preliminary homogenization was accomplished

for each sample by heating in a high vacuum to a temperature

about 200° C below the solidus temperature for at least three

hours. Electron microprobe data and x-ray diffraction patterns

were used to establish that the samples were satisfactorily homogenized.

A summary of the equilibration treatments for

the alloys in this system is given in Table 4.

28
Table 4

Summary of Equilibration Treatments

for Chromium- Platinum Alloys

Temperatures

(°C)

Time

Alloys

(atomic % Cr)

o O©

f— CD

H

3 hrs. 10, 20, 30

1500° 3 hrs. 40, 50, 60

o oO

r—

H

12 hrs. 75, 80, 90, 95

1300° 2 days 10, 20, 30, 40, 50, 60, 65

70, 75, 85, 90

1200° 3 hrs. 30, 40, 50, 60, 65

1100° 30 days 30, 76, 78, 80, 82, 84, 86

0 rH oo

r—H 14 days 75, 90

980° 17 hrs. 20, 50, 60, 65

970° 38 days 65

950° 27 days 18

910° 3 hrs. 10

900° 3 hrs. 10

900° 42 days 65, 75, 80, 90

850° 3 hrs. 18

800° 1i day 30, 40, 50

800° 60 days 50, 60, 70, 85

780° 3 hrs. 10

700° 3 hrs. 10

600° 3 days 10

29
4. Platinum Terminal Solid Solution - v and y

1

It has been established in previous studies that

chromium will dissolve in face-centered cubic platinum to form

(2- 51

an extensive solid solution region. v ’ It has also been established that atomic ordering occurs within this solid solution when

the alloys are slowly cooled or annealed at temperatures below

about 1200° C. ^ The present study confirms these conclusions,

Atomic ordering has been observed in alloys containing from

about 20 to 65 atomic percent Cr. This has been established by

the observation of superstructure lines in our Debye-Scherrer

x-ray diffraction patterns,. At the stoichiometric composition,

(6 S 8) Pt^Cr, the atomic ordering produces a CugAu(Ll2) type structure,

and complete atomic ordering may apparently exist.

^

At the

composition PtCr, the atomic ordering produces a CuAu(Ll 0 ) type

structure ^but the degree of ordering has not yet been determined*

For other compositions within the ordered region, there is apparently a continuous variation in site occupancies so that the structure

changes progressively from the Cu^Au type to the CuAu type with

fO\ *5 increasing chromium content/ ; The atomic ordering reaction

takes place quite rapidly, particularly in alloys containing over

20 atomic percent Cr, for which it is usually not possible to retain

the completely disordered condition in alloys which have been

rapidly quenched from above the order-disorder transition temperature. The temperatures of the order-disorder transition have

been reported for several alloys containing less than 50 atomic

percent Cr and are plotted in Fig. 4, (4,6) The approximate locations of the order-disorder transition temperatures for alloys

containing 18 and 20 atomic percent Cr were established by quenching sealed quartz tubes containing the alloy filings into cold water,

and noting from the x-ray patterns whether or not the alloys were

ordered. These points are also shown in Fig. 40 A smooth curve

may be drawn through these points in order to roughly delineate

the region of the ordered y

l phase, and it appears reasonable to

extrapolate this curve until it intersects the solvus line. The solvus

30
line was established mainly by the use of electron probe microanalysis of alloys quenched from various temperatures within the

two-phase y + g or y' +3 regions 0 A sharp change in the direction of the solvus line apparently exists at about 970° C, and this

sharp change in solubility may well be associated with the orderdisorder transition since it seems to occur approximately at the

point where the extrapolated order-disorder transition temperature

curve would intersect the solvus line. Further data is needed to

verify this suggestion and such data could perhaps be obtained

through high- temperature x-ray diffraction studies.

Attempts were made to determine the orderdisorder transition temperatures by the use of high- temperature

dilatometry methods 0 Dilatometric measurements were conducted

on a cylindrical sample of CrggPtgQ, which had been previously

given a homogenization annealing treatment. Filings from this

alloy had yielded an x-ray diffraction pattern of a disordered

face-centered cubic structure when these filings were quenched

from 980° C. The cylindrical dilatometer sample was, therefore,

heated to 1000° C in a high vacuum in hopes of seeing an inflection

in the dilatometer curve somewhere between 900° C and 1000° C,

which could be associated with the order-disorder transition.

However, the data revealed a smooth curve from 30G°C to 1000°

C

with no inflections. This data was kindly obtained for us by the

Argonne National Laboratory.

A similar attempt was made to detect the orderdisorder transition in a homogenized cylindrical sample of Cr^Pt^

which was heated and cooled in a high temperature vacuum dilatometer between 25°C and 1300°C.

^

Again, the data indicated a

smooth curve with no inflections, either on heating or on cooling.

The sensitivity of this dilatometer was believed to be sufficient to

detect a volumetric change in this sample of less than 0.1%. This

data was obtained by Mr. Joseph Valentich at the Westinghouse

Research Laboratories.

31
Ordering reactions, such as that which is observed in the alloy Cr

50

Pt

5() , are usually accompanied by a change

in crystal symmetry together with increases in hardness and flow

stress as the alloy transforms from the disordered face-centered

cubic to an ordered tetragonal CuAu type structure. This

reaction is accompanied by the accumulation of sizable internal

strains arising from the lattice misfit between mutually coherent

parent and product phases. The transforming system frequently

attempts to relieve these strains through shear and twinning of

the ordered phase. All of the previously known structures of the

CuAu type are in fact tetragonal.

^

However , x-ray diffraction

patterns from filings of the alloy Cr^Pt^ indicated that even

though this alloy possesses the CuAu type structure, it is apparently still cubic. High angle diffraction lines show well-resolved

doublets but no detectable splitting. Nevertheless, one may observe

profuse twinning in the microstructure of this alloy as shown in

Fig. 5a. The factors responsible for this twinning are not known

but it appears that it is not brought about by a charge in crystal

symmetry. Variations in lattice parameters are shown in Fig. 6.

5. Chromium Terminal Solid Solution - «

The body-centered cubic chromium solid solution

can dissolve a maximum of about 10 atomic percent platinum at

1500° C. However, at lower temperatures there is a significant

decrease in solubility so that near 800° C the maximum solubility

is about one atomic percent platinum. Chromium-rich alloys,

which have been annealed at low temperatures, frequently contain

a fine precipitate of the 3 phase. This fine precipitate is responsible for our inability to obtain an electron probe microanalysis

of the chromium-rich phase in alloys annealed at 900° C, since

it was virtually impossible to focus the beam into a single -phase

region. Precipitation effects are also visible in our x-ray diffraction patterns, where one observes some sharp lines of the 3 phase

together with rather diffuse lines of the body-centered cubic

32
160X (a)

50 at. % Cr, 50 at. % Pt equilibrated at 800° C for two

months. The appearance of microtwinning is difficult

to explain since x-ray diffraction data indicates that

this alloy possesses a cubic ordered structure of the

CuAu type. There was no detectable evidence of

tetragonality.

160X (b)

60 at. % Cr, 40 at. % Pt equilibrated at 800° C for two

months. The fine structure consists of p phase particles in a matrix of y'. The fine structure has nucleated at a prior grain boundary of y phase and is now

advancing toward the interior of the large y phase grain.

Fig. 5 - Microstructures of chromium-platinum alloys.

33
LATTICE

PARAMETER

0 10 20 30 40 50 60 70

ATOMIC % CHROMIUM

Fig. 6 - Variation in the lattice parameters of Cr-Pt alloys as a function of alloy

composition. Comparison is made with the data of Raub and Mahler

(Ref. 5).

34
160X (a)

65 at. % Cr, 35 at. % Pt equilibrated at 980° C for 17

hrs. Large grains are y phase containing etch pits and

inclusions of Cr2 03 . Grain boundaries are free of the

0 phase.

8 OX (b)

65 at. % Cr, 35 at. % Pt equilibrated at 950° C for 27

days. A small amount of (p + y

# ) now appears at

grain boundaries of the y phase.

Fig. 7 - Microstructures of chromium -platinum alloys.

35
160X (a)

65 at. % Cr, 35 at. % Pt equilibrated at 900° C for 42

days. The fine structure consists of 0 phase particles

in a matrix of y

' and has entirely consumed the large

prior grains of y phase. Dark spots are oxide inclusions.

370X (b) 70 at. % Cr, 30 at. % Pt equilibrated at 800° C for

four months. A prior structure consisted of 0 phase

islands in a matrix of y phase and was formed during

annealing at 1300° C. The former y phase has transformed to 0 + y'. Clear areas are prior 0 phase islands.

Fig. 8 - Microstructures of chromium -platinum alloys.

36
160X (a)

75 at. % Cr, 25 at. % Pt equilibrated at 1400° C for 12

hrs. Islands of y phase in a matrix of 0 phase.

208X (b)

76 at. % Cr, 24 % Pt after arc -melting. Dendritic 0

phase and interdentritic (0 + y) eutectic structure.

Fig. 9 - Microstructures of chromium -platinum alloys.

37
2Q8X (a)

88 at. % Cr, 12 at. % Pt after arc -melting. Area on the

left side was in close proximity to the cold copper crucible

and the cooling rate is most rapid here. Dark islands in

this region are an unresolved fine eutectic structure of

(a + 0) which tends to form a coarser (a + 0) structure in

those regions which have cooled more slowly. See Fig. 11b.

2500X (b)

The dark islands appearing in Figs. 10a and lib are

shown under high magnification. The structure is apparently that of a fine eutectic consisting of lamellar

0 phase in a matrix of a phase.

Fig. 10 - Microstructures of chromium -platinum alloys.

38
i

fc^wVSwMS!»- cM&yy^ ?s|§N ' g

pp$ -

SOX (a)

86 at. % Cr, 14 at. % Pt after arc -melting. This region

has cooled at a slower rate than the region shown in

Fig. 10b and contains dendritic p phase with interdendritic

partially decomposed a phase.

160X (b)

80 at. % Cr, 20 at. % Pt after arc -melting. Dendritic

P phase (grey) with interdendritic regions containing

two different eutectic structures. Dark islands are

the Cr-rich eutectic of (a+ p) shown in Fig. 10a. Lighter

etching fine structure is the Pt-rich eutectic of (y+ p)

shown in Fig. 9b. The x-ray pattern of this alloy contains mainly the lines of the p phase but also contains

two distinct lines identified as the (111) of P phase and

the (110) of a phase.

Fig. 11 - Microstructures of chromium -platinum alloys.

39
f ' r7^ " V , ; - - » \> * V

' .?

c

. >S°c>0 11 v' ’ ° ,v x‘ O °

: - \ \ Kd

v

-

•

.. yx O '• Oq ^0>o n«.

^ '-

s X C>'- V -X m

V </ sX~&‘S «<,„

^s.

•

• «JW&e

-.. ^ v

* /7_ t? cfcO

' c v.a;.. « > • " o 0

V o O O'

4> Xt * • r->G

\ - e^v ff ^ % 'r-: '

.. v o

• v -

l?

“ ».

° .M O t~. £ - ' t Mto o D O

• 0^* CiC -?K \\v* nh:"CX t>J

;

V'li

[X..

,• . \

c

. ( v

.. - X* • »

v*x

6

f #m d (, O

i ov

13 OX (a)

84 at. % Cr, 16 at. % Pt equilibrated at 1100° C for one

month. Globular a phase in a matrix of p phase. Grey

inclusions are Cr2 03

.

160X (b)

90 at. % Cr, 10 at. % Pt equilibrated at 1400° C for 12

hrs. Islands of a phase containing some precipitated

P phase in a matrix of p. The cooling rate after

quenching in this region was not sufficient to prevent

precipitation in the supersaturated a phase.

Fig. 12 - Microstructures of chromium -platinum alloys.

40
structure in powders annealed at low temperatures or slowly

cooled. Our cooling rates during quenching of the chromiumrich alloy powders were apparently always too slow to prevent

some decomposition of the chromium solid solution which exists

at high temperatures.

6. Intermediate Phase, 3

(5) Raub and Mahler v ' were the first investigators

to report the existence of an intermediate phase having the A15

( (3-W) crystal structure in the chromium-platinum system. They

suggested that this phase possesses the stoichiometric composition

PiCr 3 , and that it could not have a composition range since its

lattice parameter did not change with composition. We have observed that the lattice parameter of this phase does indeed show

a composition dependence, and that it varies from an = 4. 684 to

aG

= 4.708 with increasing platinum content/ Quantitative

electron-probe microanalyses of equilibrated two-phase alloys

has confirmed a previous report that the composition range of

( 13) this phase does not include the stoichiometric composition PtCrg' '

.

The maximum solubility range of this phase occurs at about 970°C,

where it is stable from 76. 5 atomic percent Cr to 83. G atomic

percent Cr.

7. Syntectic: + L2

<=* 3

Raub and Mahler^ reported that eutectic

formation occurs on both sides of the 3 phase, while the phase

itself forms under a maximum in the melting curve. The present

study is in agreement with these observations, but our data

indicate that the 3 phase may, in fact, decompose at high temperatures via a syntectic reaction rather than by congruent melting

(7)

as depicted in Fig. 321 of Hansen and Anderko’s review/ ’ We

have arrived at this conclusion after a careful study of alloys in

the "as-cast" condition which were prepared in 2 % increments

over the composition range from 76 to 86 atomic percent chromium.

41
These alloys were arc melted and then slowly solidified by gradually

reducing the arc current. This procedure produces a large temperature gradient between the bottom of the alloy sample (which is in

contact with the copper hearth) and the top of the alloy sample,

which is the last portion to solidify. This large temperature

gradient is apparently responsible for significant variations in

the microstructure of each alloy parallel to the direction of the

temperature gradient. Certain characteristic microstructures

have been observed, however, which are identifiable as eutectic

structures. The microstructural identification is confirmed, in

most cases, by x-ray diffraction patterns obtained from powders

of these "as-cast” alloys.

It appears to be quite significant that in none of

our "as -cast" alloys were we able to obtain more than about 70%

of the' 3 phase. Thus, if a congruent melting point were to exist,

then it might be expected that at least one of our alloy samples

would be in close enough proximity to the composition of congruent

melting that almost all of this sample would crystallize as the 3

phase. On the contrary, both x-ray diffraction and metaliographic

studies revealed the presence of appreciable amounts of either the

Cr solid solution or the Pt solid solution, together with the 3

phase in all of the "as-cast" samples. For alloys in the vicinity

of 80 atomic percent Cr, one observes both the Cr solid solution

and the Pt solid solution present simultaneously, together with

the 3 phase. This is rather remarkable, especially since the

microstructure of "as-cast" Crg^PtgQ (Fig. 11b) contains two

distinct eutectic structures coexisting in this alloy with the 3

phase. A photomicrograph of the dark-etching area in Fig. 11b

is shown in Fig. 10b in order to verify that this area is in fact a

eutectic structure. This dark-etching structure has been associated with a Cr-rich eutectic existing at about 85 atomic percent

chromium, whereas the lighter etching eutectic is associated

with a Pt-rich eutectic existing at about 76 atomic percent

chromium. These relationships are supported by our metaliographic and x-ray diffraction data. (See Figs. 9b, 10a, 10b).

42
It is difficult to explain the coexistence of two different eutectic

structures in a binary "as-cast" alloy. The common eutectic or

peritectic reactions do not produce such a structure. Our observation that the alloy Crg

2 Ptjg melts at about 1580°C seems to

rule out a monotectic reaction. Therefore, we are led to conclude

that only a syntectic reaction could produce the observed coexistence

of both a Pt-rich and a Cr-rich eutectic structure in our "as-cast"

alloys.

8 . Eutectic: L v + 3

Fig. 9b shows this eutectic structure in the "ascast" alloy Cr^gPt

2

^. X-ray diffraction patterns from powdered

samples of this alloy indicate the presence of v phase and 3 phase

only. Melting point determinations indicated a eutectic temperature of about 1530°C. Decreasing amounts of this eutectic structure were obtained in alloys containing higher chromium contents.

9o Eutectic: L <=* « + 3

Fig. 10a shows the structure of this eutectic in

an "as-cast" alloy, CrggPt^. The type of structure observed

seems to depend rather strongly on the cooling rate since a wide

variety of structures were observed in the "as-cast” alloy. X-ray

diffraction patterns of the "as-cast” alloy indicated only « and 3

phases, while melting point studies suggested a eutectic temperature of about 1500° C. Decreasing amounts of this eutectic structure were observed in other "as-cast" alloys in the vicinity of

this composition.

43
REFERENCES

L Muller, L. , Determination of the Melting Points of Platinum

Alloys, Ann0 Physik (5), 7, ppD 9-47, (1930).

2 0 Nemilow, W. A. , The Alloys of Platinum with Chromium,

Z. Anorg0 Chem0 , 218, pp. 33-44, (1934) 0

3. Friederich, E. and Kussmann, A., The Ferromagnetism of

Platinum-Chromium Alloys, Phys. Zeit. , 36, pp Q 185-192,

(1935).

4. Gebhardt, E. and Koster, W. , Study of the Constitution of

the Platinum- Chromium System, Z. Met., 32, pp. 262-264,

(1940).

5 0 Raub, E. and Mahler, W. , The Alloys of Chromium with

Platinum, Iridium, Rhodium and Ruthenium, Z D Met., 46,

pp. 210-215, (1955).

6o Kussmann, A., Muller, K. and Raub, E., Thermomagnetic

Measurements on Alloys of the Platinum Metals with Chromium,

Z. Met. , 59, pp. 859-863, (1968).

7 0 Hansen, M0 and Anderko, K. , ’’Constitution of Binary Alloys”,

McGraw-Hill, New York, (1958).

8 0 Pickart, S. J. and Nathans, R. , Neutron Diffraction Investigation of Pt- Based Alloys of the First Transition Series, J. Appl.

Phys ., 34, pp. 1203-1204, (1963).

9. Valentich, J. , Thermal Expansion Measurement, Instruments

& Control Systems, 42 (Oct), pp. 91-94, (1969).

10o Arunachalam, V. S. and Cahn, R. W. , Order-Hardening

in CuAu, J. Mat. Science , 2, ppa 160-170, (1967).

11. Laves, F. , Crystal Structure and Atomic Size — Table XXXI,

p. 186, in: ’’ Theory of Alloy Phases”, Trans. ASM, 48A,

American Society for Metals, Cleveland, Ohio, (1956).

44
12 0 Waterstrat, R. M. and van Reuth, E. C., Effects of Compositional Variations on the Atomic Ordering in A15 Phases,

Proc. Third Bolton Landing Conf. on Ordered Alloys,

Claitor's Publishing Division, Baton Rouge, La., (1970).

13. Waterstrat, R. M. and van Reuth, E. C., Nonstoichiometric

A15-Type Phases in the Systems Cr-Pt and Cr-Os, Trans.

Met. Soc. AIME, 236, pp. 1232-1233, (1966).

45
IVo VANADIUM-PLATINUM CONSTITUTION DIAGRAM

A0 Previous Studies

There have apparently been no previous studies designed to elucidate the constitution diagram of the vanadiumplatinum alloy systemc However, there has been sufficient

study of individual alloys to establish that various intermediate

phases are formed and to determine their crystal structures.^

- ^

Some of the phases which have been identified are apparently not

in a state of stable equilibrium at room temperature/ ’ Other

phases seem to result from contamination of the alloy samples

by oxygen or by nitrogen.' ' There is some evidence that a meta-

(41 stable phase can be obtained by cold- working the alloy powders/ ’ An investigation of the constitution diagram is, therefore, essential

in order to identify which of these structures belong to equilibrium

phases, and perhaps also to suggest some reasons for the existence

of the various structures which have been observed.

B. Materials

The vanadium metal used in this investigation was

obtained from the UQ S. Bureau of Mines, Boulder City Metallurgy

Research Laboratory through the courtesy of Mr. T. A. Sullivam

The metal was in the form of electrolytic dendrites having a nominal purity of about 99 c 95%. A typical analysis of this metal is

shown in Table L

The platinum metal was in the form of l/32-inch sheet

and was from a stock supply whose original source was unknown.

Spec tro graphic analysis, shown in Table 5, revealed that this

metal had a nominal purity of about 99.9% with the principal

impurity being palladium which was present to the extent of

between 100 and 1000 ppm.

46
Table 5

Typical Analyses of the Materials Used in

Preparing the Experimental Alloys +

Impurity

Vanadium Platinum

U. S. Bureau of Mines Stock Supply

Ag ND <10

C 80 -- Cr 60 ND

1 Cu 32 <100

Fe 135 <10

H 10 -- Mg ND <100

N 35 -- 0 220 --

Pd ND <1000

Ru ND <100

Si ND <100

t Values are given in ppm.

ND Not detected spectrographically.

No other elements were detected spectrographically.

C. Alloy Preparation

The experimental alloys were prepared by weighing

the desired amount of each constituent metal and then melting

these metals in an arc furnace under an inert atmosphere.

47
1. Weighing

The components of each alloy were weighed to

an accuracy of ± 0. 01 gram, and each sample had a total weight

of about 20 grams. The weights were checked after each melting

in order to establish that no excessive weight changes had occurred.

In most cases, there was very little weight change, and in all

cases weight changes were less than 1%.

2 C Melting

All samples were arc melted on a water-cooled

copper hearth using a non-consumable tungsten electrode and a

furnace atmosphere of special high purity helium and argon in

equal proportions by volume. The purity of these gasses is given

in Section III, Table 2. Techniques of melting were similar to

those which are described in Section III, C-2. Alloys rich in

vanadium were melted using cylinders of compressed vanadium

dendrites which were prepared by placing the appropriate quantity

of vanadium in a 1/2-inch diameter hardened-steel die and compressing this in a hydraulic press until a firm, compact cylinder

was obtained. In some cases, the platinum sheet was wrapped

around these cylinders prior to melting the alloys.

Special care was taken with brittle alloys to

insure that no losses of metal occurred when these alloys shattered

under the impact of arc initiation. Such losses could usually be

prevented by breaking up the alloy before melting. In cases where

such shattering still occurred, we exercised scrupulous care in

collecting all scattered fragments for re melting.

D. Composition Determination of the Melts

Since weight changes during alloy preparation were

always quite small, it appeared safe to assume that the actual

compositions of the alloy samples were quite close to the nominal

48
intended compositions. This assumption was verified throughout

the investigation by the results of electron probe microanalyses 0 Our sample, V^Pt^ (nominal), was submitted to a chemical

analysis which indicated that its composition was indeed 50.0

atomic percent V and 50. 0 atomic percent Pt. A portion of this

same sample, together with samples of pure platinum and pure

vanadium, were then used for calibrating the electron probe

data on a quantitative basis.

Lattice parameter measurements were of little value

in determining alloy compositions since they were rather insensitive to composition apparently due to the close similarity in the

atomic diameters of vanadium and platinum.

E. Temperature Measurement

A platinum versus platinum- 10% rhodium thermocouple

was used in measuring temperatures up to 1500° C. Temperatures

above 1500°C were measured with a tungsten versus tungsten-26%

rhenium thermocouple. These thermocouples were periodically

calibrated against a standard platinum-6% rhodium versus platinum30% rhodium thermocouple which had been calibrated at the

National Bureau of Standards up to 1600° C. The thermocouples

were also calibrated at secondary fixed points, such as the melting

points of nickel or platinum, using the methods described in

connection with our solidus determinations.

F. Thermal Treatments

L Homogenization

All samples were homogenized by heating to

within about 200° C of the solidus temperature for a length of time

which could be considered sufficient to establish chemical equilibrium at that temperature. Thus, the samples could be considered

49
as being equilibrated at the homogenization temperature since

all precautions normally used in an equilibration treatment were

observed.

The samples were suspended in a tantalum

bucket within the hot zone of our tantalum-strip resistance furnace 5 operating at a pressure of between 5 x 10~® and 5 x 10”^

mm Hg. Any reaction with the tantalum bucket was apparently

insignificant, and was not detected by the electron probe. The

electron probe also indicated that our samples had been adequately homogenized since no composition gradients were detected.

This was further verified by metallographic and x-ray diffraction

studies.

2. Equilibration

Most of our annealing treatments at temperatures up to 1100° C were done by sealing the samples in quartz

tubes which had been baked- out, evacuated, and filled with a

partial pressure of high purity argon. However, alloys containing between 33 and 60 atomic percent vanadium were found to be

seriously contaminated at 1100° C, apparently by a reaction with

the quartz tubes. An attempt was made to prevent this reaction

by wrapping the samples in tantalum foil, but this was unsuccessful. Many small crystals were observed on the interior walls

of the tubes. These were collected and submitted to x-ray diffraction and electron probe analysis. The crystals were found to

contain vanadium, tantalum, and silicon. X-ray precession

pictures indicated that the crystals were tetragonal with possible

space groups I4cm, 14 c2, or 14/mem. The latter space group

is that of the compound VgSig. The lattice constants of our

crystals (a0

= 6.46, c 0 = 12.08), however, resemble those of

NbgSig (aQ

= 6. 57, c Q

= 11.88). We concluded that a VanadiumTantalum Silicide had been formed.

50
In order to prevent this contamination, we decided to seal our samples in pure platinum tubes, which were then

evacuated and filled with an appropriate partial pressure of pure

argon before closing the tubes by compressing the ends. This

method seemed to be quite satisfactory in preventing contamination of the alloy samples. It was essential that the tubes contain

the correct amount of argon gas so that they would not collapse

when heated to the annealing temperature. A platinum tube, which

collapses on the samples, will often rupture in the vicinity of any

sharp edges on the samples.

All samples were suspended in a vertical

platinum- wound resistance furnace using either nichrome or

platinum wire to hold the tubes. Some of these samples were

also annealed in our tantalum- strip resistance furnace in a high

vacuum, and no contamination was detected under these conditions. The tantalum- strip resistance furnace was also used for

all annealing treatments above 1100° C. Pressures of between

5 x 10" ^ and 5 x 10“ ^ mm Hg were obtained during the annealing

treatments and temperatures were controlled to within ± 10° C„

All samples were rapidly cooled from the

annealing temperature by turning off the furnace power. The

cooling rate was sufficiently rapid to permit identification of the

prior structures existing at the annealing temperature, either

by metallo graphic or x-ray diffraction studies. Samples annealed

in quartz tubes were quenched in cold water. A summary of the

equilibration treatments for the vanadium-platinum alloys is

given in Table 0 O

G. Methods for Determination of Phase Boundaries

I. Solidus Temperatures

The procedures used in determining the solidus

temperatures were similar to those used in our study of the

51
chromium- platinum system. (See Section III, G-h) Although

there is no problem of high vapor pressures, as was the case

with the chromium alloys, there is nevertheless a problem in

preventing chemical reactions between vanadium and some refractory materials at high temperatures. In most cases, this

was solved by using thoria crucibles in our solidus measurements, but high purity alumina crucibles were satisfactory for

the platinum- rich alloys.

Our success in determining solidus temperatures by direct observations of melting for the chromiumplatinum alloys encouraged us to use similar methods for the

vanadium- platinum alloys. The method appeared to be equally

satisfactory in this case, partly because a congruent melting

maximum exists in both the Cr-Pt and V-Pt systems, and this

gives rise to sharply defined melting points. The method is

described in Section HI, G-l, and illustrated in Section III,

Fig. 2 C No special modifications were required for the

vanadium- platinum alloys other than the use of thoria crucibles

for alloys containing more than 30% vanadium. Solidus temperatures, as much as 50° lower, were obtained in some cases when

alumina crucibles were used with these alloys. The thoria

crucibles appeared to be quite resistant to any reaction with

vanadium, and for this reason we feel that our solidus temperatures are fairly reliable.

2. Liquidus Temperatures

In most cases, we did not attempt to determine

the liquidus temperatures, but in the region of the 3 phase, we

were able to observe an appreciable temperature interval between

the first observation of melting and the temperature at which the

sample became completely molten. In these cases, we have presented the temperatures of complete melting as points on the

liquidus curve. These points should only be regarded as

52
approximate, however, and would require verification by more

accurate methods.

3. Invariant Reactions

The existence of invariant reactions and the

temperatures at which these reactions occur were 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 be retained to room temperature by rapid quenching

in each case, except for the alloy V^Pt^. In this case, high

temperature dilatometry was employed to locate the transformation temperature.

Although the comparative method may appear

to be excessively tedious, it has several significant advantages.

Perhaps the most important advantage is the fact that one can

anneal the sample for as long as is necessary to insure that an

equilibrium structure has been formed at the annealing temperature. This avoids the uncertainty of possible metastable conditions developing, as may occur when using methods employing

slow cooling. In addition, the comparative method in many cases

offers clear, unequivocal evidence as to whether or not the

sample was above or below the transformation temperature,

when with other methods the evidence may be ambiguous. Finally,

the comparative method permits one to begin with any arbitrary

annealing temperatures. Once it is established whether this

temperature is above or below the transition point, it is only

necessary to select some other annealing temperature on the

opposite side of this transition point. When the transition temperature has been bracketed in this manner, it is not very difficult to

locate the exact transition temperature since the data converge

rapidly. In most cases, less than six annealing treatments were

needed to establish the transition temperatures within the accuracy

of the temperature measurement itself.

53
Eutectic temperatures are readiiy obtained by

direct visual observations of melting using the methods described in Section IV, G-l for the determination of solidus temperatures. Metallographic examination of the melted alloys can

reveal evidence regarding the proximity of the eutectic composition Thus, metallographic examination of a few melted alloys

usually enables one to locate the eutectic composition.

4. Other Phase Boundaries

The solvus lines and the phase boundaries of

intermediate phases were obtained by electron probe microanalysis whenever this method was feasible. The procedures

are described in the Appendix, Unfortunately, the alloys in this

system were frequently unsuited to this method due to the presence

of unresolvabie fine structures in the alloy samples. These structures usually resulted from the eutectoid decomposition of phases

occurring at high temperatures. At least three separate eutectoid

transformations occur in this system, and we were therefore

forced to rely on metallographic studies to locate many of these

phase boundaries. The metallographic studies require more

effort and time than the electron probe measurements, but the

results seem to be sufficiently reliable to establish the location

of the phase boundaries to the desired accuracy. Electron probe

microanalysis was effective mainly in establishing the boundaries

of the 3 phase region in this alloy system.

H. Experimental Results

1. General

The proposed constitution diagram for the

vanadium- platinum system is presented in Fig, 13. Discussions

of various phase regions are given below. All alloys were examined either metallographically or by means of x-ray powder

54
WEIGHT

PERCENT

VANADIUM

Jo‘3UniVa3dlAI31 !II oo 8o

<T> O

CO8o

COo

IT) co

rOO

CMoo o o O o o o o o o o o o O o o o o o o o

ID ro cn h- lO rO cn IsCM CM CM — — — — — >

Cl<z<>

LlJ

C_) crUJ

Cluo

<c0.

Do‘3aniVH3dW31 3u

bD3• pH 32GO

COGOo33

333

r*3>

cu JZH

co

th

•pH

55
patterns, or both. There was generally good agreement among

the metaliographic, x-ray, and electron probe data.

2. Solidus

The solidus was determined mainly by the

technique of direct observation of melting as described in Section

III, G-l. Melting generally occurred quite rapidly at a definite

temperature except for alloys containing over 70 atomic percent

vanadium which melted over a significant temperature range.

Each solidus measurement was accompanied by a simultaneous

calibration using high purity platinum wire as a melting point

standard. The calibration was accomplished using the same

experimental conditions as was used in determining the solidus

temperature of the alloy sample. It was possible to detect fairly

small temperature differences between the melting point of the

standard wire and the temperature at which the alloy sample

began to melt since the thermocouple had a sensitivity of about

± 1°C. It was necessary to use thoria crucibles for all alloys

containing more than 30% vanadium, since vanadium apparently

reacts with alumina crucibles at these high temperatures.

3. Liquidus

In most cases, we did not attempt to determine

liquidus temperatures, but in the case of the alloys V^Ptgg,

V83 Pti7

, and V87 Pt

13 , which melt over an appreciable temperature

range, we assumed that the temperatures at which melting appeared to be completed (e.g. , the sample abruptly collapsed and dropped

into the crucible) were in fact the approximate liquidus temperatures

as shown in Fig. 13.

56
4. Equilibration

Each sample was first homogenized by annealing

in a high vacuum at a temperature within 200° C below the solidus

temperature for various periods of time. Electron microprobe

data indicated that single-phase samples were homogeneous and

that our two-phase alloys contained homogeneous phases. A

summary of the equilibration treatments is given in Table 6.

57
Table 6

Summary of Equilibration Treatments

for Vanadium- Platinum Alloys

Temperatures HPi mp Alloys

(°C) (atomic % V)

1700° 2 hrs. 10,22,23,36.38.75,92

1600°

1550°

4 hrs.

2 hrs.

46,49,50,51,53,55,57,63,87,95

50

24 hrs.

10,20,30,40,46,49,50,51,53,55,57,

1500° 61,63,65,67,69,71,73,75,77,79,81,

83,90,95

1450° 5 hrs. 59

1430° 4 hrs. 55

1430° 26 hrs. 43,46,53

1420° 4 hrs. 55

1400° 4 hrs. 55

1300° 3 days 10,20,30,38,40,43,46,48,49,51,53,

55,59,65,75,83,87,95

1200° 1/2- hr. 43

1110° 1 hr. 33

1100° 2 hrs 0 33,48

1100° 18 days 10,20,25,33,36,38,40,43,46,48,50

57,87,92,95

1090° 2 hrs. 33

1080° 2 hrs. 33

1070° 2 hrs. 30

1050° 2 hrs. 33

1030° 3 wks. 30,36,38,43,55,65,75,83,95

1020° 2 hrs. 23.5,25,26.5,28,33,40

1010° 2 hrs. 25

1000° 2 hrs. 25,36,38,40,43,46,48

995° 2 hrs. 28

990° 2 hrs. 26.5

980° 2 hrs. 22 23.5,25,28

975° 2 hrs. 38

960° 2 hrs. 28

950° 2 hrs. 20,25,30,40

940° 2 hrs. 22,28

900° 6 wks 0 10,20,22,23.5,25,26.5,28,30,33,

36,43,48,53,59,65,75,83,90,95

800° 7 days 10,20,22,23.5,25,26.5,28,30

58
5. Platinum Terminal Solid Solution - v

It was found that vanadium will dissolve in facecentered cubic platinum to form an extensive solid solution at

temperatures above 1500° C, but at lower temperatures there are

at least three different order-disorder type transformations,.

These transformations occur at the stoichiometric compositions

VPtg, VPtg, and V Pt. Each of these three order-disorder transformations is accompanied by a change in crystal symmetry0 Consequently, as the ordered phase forms in the disordered matrix,

there is an accumulation of internal strains which presumably

gives rise to the microscopical observations of shear and twinning

in the ordered phases.

In addition to these three order-disorder reactions, we have also observed atomic ordering in our alloy,

^22^78° case > there is no change in crystal symmetry

since the ordered structure is a CugAufL^) type cubic structure.

However, the fact that this ordered structure is not observed at

the ideal (ABg) stoichiometric composition, and is only found in

a region where one would expect to observe two other phases ( y +

0), raises some doubts as to whether it is in fact an equilibrium

phase in this binary system. Maldonado and Shubert v ' have

reported the formation of a Cu^Au type structure in their alloy,

V Ptg, but our alloy V Ptg did not produce this structure, even

after a similar heat treatment (5 minutes at 1000° C). Our alloy

^23 5^*76 5’ however, did exhibit the Cu^Au type structure after

annealing for two hours at 980° C, but this same alloy showed only

the TiAlg type structure of the 0 phase when annealed at 900°

C

and at 800° C. We have, therefore, designated the Cu^Au type

phase as y

1 in our constitution diagram (Fig0 13), and its approximate stability range is delineated with a dashed line 0 Further

work is needed to establish its mode of formation and whether it

is actually a binary equilibrium phase.

59
Lattice parameters were obtained for alloys

within the region of the disordered platinum solid solution, y , whenever it was possible to retain this structure by quenching

the x-ray diffraction powders from high temperatures. This

data is shown in Table 7, and indicates that increasing amounts

of vanadium in the face-centered cubic v solid solution produce

a gradual decrease in the lattice parameter. It was not possible

to retain the high temperature face-centered cubic structure by

rapid cooling of alloy powders containing more than about 40 atomic

percent vanadium. We attempted to retain the high temperature

structures by rapid cooling of the bulk samples, but this method

was found to be unreliable since it is still necessary to remove

the x-ray powders from such samples by some process whereby

the powder particles are severely deformed. There is evidence

(4) that such deformation can produce a structure change in ¥ Pt v ’ and we have observed similar behavior in V Ptg and VPtg alloys.

If one attempts to relieve the strains in these powders by reannealing them at the original annealing temperatures, then one is again

faced with the problem of obtaining sufficiently rapid cooling of

the powders. We were, therefore, unable to obtain direct evidence

of a face-centered cubic structure in alloys containing over 40

atomic percent vanadium, but the continuity of the platinum solid

solution in this region is inferred from metallographic examination

of these alloys and by other evidence which will be considered in

the subsequent discussion.

6 . Vanadium Terminal Solid Solution -«

The extent of the vanadium terminal solid

solution was determined mainly by electron probe microanalysis

of two-phase alloys annealed and quenched from various temperatures within the °= + 3 region of the diagram (Figo 13) 0 Lattice

parameter measurements were attempted, but the alloy powders

apparently undergo either a slight contamination or a precipitation

60
Table 7

Lattice Parameters of y Phase

V-Pt Alloys

Composition

(atomic % V)

Lattice Parameter - a0

(A)

0 3.924

10 3.900

20 3.887

25 3.879

28 3.875

30 3.872

36 3.864

40 3.861

of the 3 phase, which produces line-broadening effects and consequent inaccuracies in locating the positions of the diffraction

lines.

Our data indicate that platinum will dissolve

in body-centered cubic vanadium to form a solid solution which

extends to about 12 atomic percent platinum at 1800° C, but this

solubility decreases at lower temperatures. Metallographic

studies indicated that no precipitation occurs in the alloy Vq9 Ptg

if it is rapidly cooled from 1700°C, but some precipitation was

observed when this alloy was reannealed at lower temperatures.

The "as-cast” alloy appeared to possess a greater resistance

to tarnishing in air than a similar sample of pure vanadium.

61
Intermediate Phase 0

Dwight, Downey and Conner

^

have identified

the crystal structure of V Ptg as being isotypic with TiAl^, and

this has been confirmed by Maldonado and Schubert.^ The unit

cell of this structure is tetragonal and belongs to the space group

I4/mmm0 The lattice constants for V Ptg obtained in our study

were aQ = 3.839 A., c 0 = 7„796 A, and we have noted that these

values are smaller than the values reported by Dwight, Downey

and Conner.^

The experimental uncertainty in our measurements, however, is at least ± 0.01 A, and therefore the significance of this discrepancy may be questionable G Our use of very

high purity vanadium may also account for some discrepancies

with the previous data.

We have observed that the lattice constants

of the 0 phase exhibit a significant composition dependence,,

This suggests that it is stable over a composition range of

several percent. We obtained the values aQ

= 3.860, c 0

= 7.819

for the 0 phase alloy, Vg

g

gPt^g 5, while the © phase alloy,

V20

5

Pt

73 5 yielded the values a^, = 3.847, c 0 = 7.796. We have,

therefore, estimated that the composition range of the 0 phase

extends from 23 to 27 atomic percent vanadium at 800 to 900° C,

with an uncertainty of ± 1 atomic percent.

The 0 phase apparently forms from the facecentered cubic terminal solid solution of platinum via an orderdisorder reaction which occurs at 1015 ± 5°C. If the alloy V Ptg

is annealed at a temperature above this transformation temperature, and then quenched in cold water, one may retain the facecentered cubic structure to room temperature. One may also

produce the face-centered cubic structure by subjecting 9 phase

filings to severe plastic deformation at room temperature. Subsequent annealing of these severely deformed powders will apparently induce recovery and recrystallization, and one ultimately

62
obtains the equilibrium structure which exists at the annealing

temperature. Incomplete annealing of cold- worked 9 phase

filings may produce metastable structures, and this might per-

( 3 ) haps explain why Maldonado and Schubert v ' observed a Cu^Au

type structure in their alloy VPtg. However, we have never

observed a Cu^Au type structure in our alloy V Ptg for any of

the various annealing treatments used in this study, nor even for

the annealing treatment used by Maldonado and Schubert (5

minutes at 1000° C)/3 ^

As mentioned previously, however, we did

observe a CugAu type structure when we reannealed some

filings of a 0 phase alloy, Vgg gPt^g

5

? ^or two hours at 980° C.

We are not certain whether this structure represents that of a

binary equilibrium phase in this system, or whether it is stabilized by some impurity. It is interesting, however, that one

can obtain either a TiAlg type ordered structure, a Cu^Au type

ordered structure, or a disordered face-centered cubic structure in this alloy simply by annealing the alloy filings at the

appropriate temperature. The relationship of the TiAlg type

structure to the Cu^Au type structure has been discussed by

(7)

0

H. J. Beattie/ 7 and is illustrated in Figure 4 of his papera We have reproduced this illustration in our Fig. 14. Both of

these ordered structures are based on a face-centered cubic

arrangement of the atoms. The type of atomic ordering which

occurs will determine whether one obtains a cubic Cu^Au type

unit cell of approximately the same size as that of the disordered

face-centered cubic structure, or whether the atomic ordering

will essentially double this unit cell to produce a tetragonal

TiAlg type structure. It may be noteworthy that a change in

crystal symmetry occurs only in the case of the tetragonal

TiAlg type structure, and thus there will be an accumulation

of internal strains during a transformation to this structure.

These strains are relieved by shear or twinning which is

apparently responsible for the striations which are seen in Fig. 15.

63
Cu^Au - type

Fig. 14 - Relationship of the Als Ti-type and Cu3 Au-type structures.

64
Table 8

X-ray Diffraction Pattern of 6 Phase VPt^

Tetragonal TiAlg- type Structure

a

Q

= 3. 839 a, c Q

= 7. 796 a

^(obs)

U)

Relative Intensity

hkl Observed Calculated t

3.881 vw 6 002

3 0 589 vw 17 101

3.438 * ww

3. 283 * vw

2 e 716 w 5 110

2.208 vs 100 112

-- __ 5 103

1.950 m 16 004

1.922 ms 31 200

1.726 mw 2 202

1. 698 vw 4 211

1.662 * vw

1. 586 mw 2 114

1.469 vvw 3 105, 213

1.370 ms 18 204

1. 361 m 9 220

-- -- <1 006

1.247 w 1 222

1.235 vw 1 301

1.217 w 1 310

1. 175 ms 10 116

1.163 s 21 312, 215 __ -- 1 303

1.118 ms 9 224

L 079 vw 1 206,107

1.050 ww 1 321

1.034 mw 1 314

65

(Table 8 Continued)
Table 8 (Continued)

0o 9993 vw 1 305,323

0.9771 m 2 008

0.9641 ms 3 400

0 o 9416 ww <1 226

0c 9360 ww 1 217,402

-- __ 1 411

0.9207 ww <1 118

0o 9035 ww <1 330

Qo 8907 s 12 316

0o 8853 ms 6 332

-- 1 325, 413

0. 8722 ms 6 208

0c 8648 ms 6 404

0 o 8624 ms 6 420

1 109, 307

0, 8422 w 1 422

0o 8245 w <1 334

-- 1 415

0 o 7949 s 10 228

0, 7894 vs 23 424

__ — <1 0-0*10

0.7755 m 3 219, 406

m 4 327

* These very weak lines are attributed to the presence of

some unidentified impurity.

+ Relative intensities were calculated assuming complete

ordering of the vanadium and platinum atoms on the available lattice sites; two vanadium atoms in positions 000,

1/f3

l7

3 > and six platinum atoms at OQ 1 ^ , 1/2

1/2 0,

0

1

/ S 1/4 J 0V2 3/4j

1

/s 0 %, 1 / 2 0

3 U. The calculated

intensities do not include an absorption correction and

the absorption is appreciable in our sample. The observed relative intensities therefore are suitable for

comparison only among lines whose Bragg angles (or dvalues) are not greatly different.

66
I

8OX (a)

25 at. % V, 75 at. % Pt equilibrated at 900° C. Large

grains of 9 phase with microtwinning.

’

160X (b)

28 at. % V, 72 at. % Pt equilibrated at 950° C for 27

days. Fine structure is presumably due to the eutectoid decomposition of y to form 0 phase + 6 phase.

Fig. 15 - Microstructures of vanadium -platinum alloys.

67
A similar transformation occurs in the V-Ni system and is associated with age-hardening effects. An x-ray diffraction pattern of 0 phase VPt^ is presented in Table 8.

80 Intermediate phase 6

The crystal structure of the phase VPtg has

been identified as an MoPtg type structure by Schubert et al.^j^)

Their identification was confirmed in a subsequent study by

Giessen and Grant, who reported a more accurate determination of the lattice parameters. The structure is orthorhombic

and belongs to the space group Immm0 The vanadium and platinum

atoms are ordered in the two crystallographic positions. The

lattice parameters obtained in this study were aQ = 2.724 A,

b0 = 8.302 a, and c 0 = 3. 786 A, all ± 0.01 A. The values obtained

were not significantly affected by changes in the alloy composition,

which suggests that this phase is stable only within a rather narrow

composition range. Metalk) graphic studies suggested that the

composition range of this phase extends from about 32. 5 to 34.0

atomic percent vanadium at 900°C, with an uncertainty of ± 1 atomic percent.

The 6 phase forms from the face-centered

cubic platinum solid solution through an order-disorder reaction

which occurs at 1105 ± 5°C. If the alloy V Pt

2

is quenched from

a temperature above this transition temperature, one may retain

the face-centered cubic structure to room temperature. It is

also possible to produce the face-centered cubic structure by

severe plastic deformation of 6 phase alloy filings. These filings

may be subsequently annealed to produce an equilibrium structure,

and this occurs in a manner quite similar in behavior to that which

has been described for the 0 phase alloys. In both cases, it is

possible that the structure transformation is produced by the

destruction of long-range ordering resulting from severe coldworking of the alloy powders. It is also possible that incomplete

68
Table 9

X-ray Diffraction Pattern of 6 Phase V Pt2

Orthorhombic MoPt2~type Structure

a0

= 2. 724 A, b0= 8. 302 A, cQ

= 3. 786 A

d(obS ) Relative Intensity

(a) Observed Calculated +

4.107 w 16 020

3.429 m 44 011

2. 575 mw 20 110

2.230 vs 100 031

2.209 vs 99 101

2.073 w 9 040

1.950 w 8 121

1.932 ms 63 130

1.893 m 30 002

1.723 w 3 022

1. 530 mw 8 112

1. 511 m 12 141

1. 399 w 5 042

1. 382 m 9 060

1. 363 m 9 200

1.356 s 35 132

1. 292 vw 1 220

1. 267 w 4 211

1.252 vw 2 013

1. 172 ms 19 161

1. 162 ms 20 231

1. 149 m 19 033, 103

1. 139 vw 3 240,152

1. 133 w 3 071

1. 118 m 8 062

1. 107 m 10 202,123

(Table 9 Continued)

69
Table 9 (Continued)

L 089 w 3 170

L 069 w 1 222 — 0 080

lo 014 vw <1 251

1.005 mw 3 053,143

0 o 9770 w 3 242

0» 9718 m 6 260

0.9505 w 3 004

0 o 9449 mw 4 172 — -- <1 181

0 o 9222 w 2 024, 213

0 082

0„ 9043 vw 1 310

0 o 8980 m 5 091

0o 8875 vw 2 114

0 o 8844 ms 17 301, 163

0 o 8799 m 12 233

0„ 8751 w 9 190, 271

0 e 8651 s 32 073, 262, 321,330, 044

0.8541 ms 13 134

0.8312 w 2 0 *10 *0 -- -- 0 280

0. 8161 mw 2 312

0. 8140 mw 4 341

-- <1 253

0.7952 ms 17 350, 192

Go 7894 mw 1 154

0 o 7867 s 24 332

0.7799 m 13 064

0.7780 m 16 204

0.7770 m 15 1*10*1

70
t Relative intensities were calculated assuming complete

ordering of the vanadium and platinum atoms on the

available lattice sites; two vanadium atoms in positions

2(a) (000) ( 1/

2

X/a ^a) and four platinum atoms at positions

4(g) (OyO) (V 2 v2 + y,v2 ) (0, - y,0) - y,

v2 ) where

y

= 0. 341. The calculated intensities do not include an

absorption correction and this may be appreciable in our

sample. The observed relative intensities are therefore

suitable for comparing only those lines whose Bragg angles

(or d-values) are not very different,

71
160X (a)

33 at. % V, 67 at. % Pt equilibrated at 900° C. Large

grains of 6 phase heavily etched to reveal microtwinning.

400X (b)

40 at. % V, 60 at. % Pt equilibrated at 1300° C for 3

days. Alloy is composed of the y solid solution with

no evidence of any second phase.

Fig. 16 - Microstructures of vanadium -platinum alloys.

72
annealing of the cold- worked powders would produce a metastable

phase, and this may account for the observation of a CuAu type

structure in certain V-Pt alloys.^

The formation of the ordered phase V Pt

2

from

the face-centered cubic solid solution involves a change in crystal

symmetry, as in the case of VPt^, and consequently, one observes

similar evidence of the existence of internal strains. This can be

seen in the form of striations within each grain as shown in

Fig. 16a, One may also expect to observe age -hardening effects

in these alloys. An x-ray diffraction pattern of the 6 phase VPt

2

is presented in Table 9.

9. Intermediate Phase g

Schubert et al.^’^ were the first to report

the existence of a phase VPt having the AuCd-type structure.

We have confirmed this identification, but we were unable to

confirm their identification of a CuAu type structure at this

same composition, except for alloy filings which had been severely

deformed by cold-working. The severely deformed alloy powders

produce x-ray diffraction patterns having rather broad lines which

did not permit careful indexing, but nevertheless, these lines did

(4)

appear to be those of a CuAu type ordered phase.' ’ If these deformed powders were reannealed at various temperatures, however,

they would invariably transform to the AuCd type structure. Thus,

it appeared that the CuAu type structure was a metastable phase,

and that the true equilibrium structure of the phase VPt is the

AuCd type structure.

Metallographic examination of the alloy V Pt,

however, revealed profuse twinning (see Fig. 20b) and this at

first suggested to us that the AuCd type structure may have been

formed by a martensitic transformation. If such were the case,

then it would be likely that some unidentified structure existed at

73
high temperatures which had probably not been retained to room

temperature, even by rapid cooling. On the other hand, it also

appeared possible that the AuCd type structure, which is highly

ordered, may have been formed by an order-disorder transformation from an extension of the face-centered cubic platinum

solid solution in a manner similar to the mode of formation of

V Ptg and V Ptg.

/ An attempt was therefore made to detect evidence of any structure transformations which may occur in this

alloy from room temperature up to within about 100° C of its

melting point. This was done by using a high temperature vacuum

dilatometer to detect small changes in specimen length during

several heating and cooling cycles. This dilatometer had been

designed and constructed by Mr. Joseph Valentich at the Westinghouse

Research Laboratories,^ who kindly conducted the measurements

on our sample. The sample was prepared by carefully grinding

a portion of an arc melted button to obtain a cylindrical sample

of about 1/4-inch diameter by 1-1/4-inches long. The alloy is

quite brittle, and we are indebted to Mr. John Kumpula of the

National Bureau of Standards for his skillful and patient effort

in preparing this sample without any visible cracks. The cylindrical sample was then homogenized and stress-relieved by

annealing it in a tantalum V-block holder at a temperature of

1600° C, followed by slow cooling to room temperature.

Subsequent heating and cooling cycles in the

high temperature dilatometer from 25°C to 1650°C to 25°C

revealed evidence of only one transformation occurring at about

1500 ± 10° C. Data for a typical heating and cooling cycle are

shown in Fig. 17. The arrest was fairly sharp and occurred at

nearly the same temperature during heating and during cooling.

A volume change of about 0. 6% is suggested by this data with the

low temperature phase apparently having a smaller specific

volume (i. e., greater density). The dilatometer is capable of

74
V

oc<

IT

£Lm

yjH

75

Fig.

17-

Thermal

expansion

of

VPt

asa

function

of

temperature.
detecting volumetric changes of less than 0.1% and we have,

therefore, concluded that the AuCd type structure which we have

observed in our V-Pt alloys is probably stable from room temperature up to 1500°C, where it apparently transforms to another

structure which is stable only above 1500° C, and probably up to

the melting point of this alloy (-1770° C). Metallographic studies

of alloys in the equiatomic composition region suggest that the

high temperature phase is, in fact, a continuation of the facecentered cubic platinum solid solution from which the £ phase

forms by an order-disorder reaction at about 1500°C. The onset

of ordering apparently produces a distortion of the face-centered

cubic structure resulting in the orthorhombic AuCd type structure.

It is interesting to observe that the AuCd type structure has been

considered to be a distortion of a close- packed hexagonal structure, a distortion of a body- centered cubic structure, and,

as in the present case, a distortion of a face-centered cubic

structure.

It appears that the £ phase coexists with the 3

phase on the V-rich side of the equiatomic stoichiometry as a

result of a eutectoid decomposition of the high temperature y

phase. On the Pt-rich side, however, the £ phase coexists

with the face-centered cubic y phase, and apparently this gives

rise to the unusual microstructure shown in Fig. 18b. This

structure consists of parallel ribbons of the £ phase apparently

having a strong epitaxial relationship with the matrix y phase,

since this structure is stable even after annealing for three

days at 1300° C. The ribbons are not simply due to twinning of

the y phase since, first, they have a different crystal structure

which is revealed strongly by polarized light, and secondly,

they have a different composition (richer in vanadium) than the

matrix phase as revealed by an electron beam target current

scanning image shown in Fig. 19a . The formation of ribbons

of the £ phase, which are strongly coherent with the face- centered

76
208X (a)

43 at. % V, 57 at. % Pt equilibrated at 900°C for 6

weeks. Prior y phase regions have decomposed

eutectoidally to form 6 phase + Q phase.

500X (b)

43 at. % V, 57 at. % Pt equilibrated at 1300° C for 3

days. Ribbons of Q phase in a matrix of y phase. See

Fig. 19a.

Fig. 18 - Microstructures of vanadium -platinum alloys.

77
10QQX (a)

43 at. % V, 57 at. % Pt equilibrated at 1300°C for 3

days. This is an electron beam target-current scanning image of the structure shown in Fig. 18b. It reveals that V-rich ribbons (dark) co-exist with a

Pt-rich matrix (light).

SOX (b)

46 at. % V, 54 at. % Pt equilibrated at 1300°C for 3

days. Bands of VPt in a matrix of y phase.

Fig. 19 - Microstructures of vanadium -platinum alloys.

78
400X (a)

48 at. % V, 52 at. % Pt equilibrated at 1300° C for 3

days. Small globules of y phase at grain boundaries

of the Q phase.

52X (b) 50 at. % V, 50 at. % Pt equilibrated at 1500° C. Q

phase grains showing microtwinning.

Fig. 20 - Microstructures of vanadium -platinum alloys.

79
cubic v phase, suggests interesting possibilities of fibre reinforcement to obtain superior mechanical properties„ Unfortunately, alloys in this composition range seem to suffer from

some form of intergranular embrittlement, and it may therefore be necessary to first grow single crystals of the y phase

in order to take full advantage of the reinforcement provided

by the £ phase ribbons.

The AuCd type structure is orthorhombic

and belongs to the space group Pmma. We have indexed all

the lines in the powder pattern of the £ phase, and have also

calculated the relative line intensities. This is shown in Table

10, The intensity agreement is satisfactory for all lines except

(200). The intensity of the (200) line appears to be significantly

stronger than the calculations would indicate, and the reason

for this discrepancy has not yet been discovered. Attempts

to improve the intensity agreement by assuming variations in

the degree of long-range atomic ordering, or by variations in

the atomic position parameters z e or Zf, were unsuccessful.

The final lattice parameters obtained for our alloy VPt were

a0

= 4.413 a, b0= 2.693 A, c 0

= 4. 767 A ± 0.005 A. There was

apparently some change in line positions produced by composition variations, but these were rather small and we did not

attempt to determine the variation in lattice parameters. It appears, however, that the £ phase does possess a significant

compositional range of stability which we have estimated from

metallographic observations as being from 49. 0 to 51. 5 atomic

percent vanadium ± 1 atomic percent.

It was observed that vanadium- platinum alloys

having compositions in the equiatomic region are quite susceptible to contamination, particularly when the alloy lumps or

powders are heated in quartz tubes. In such cases, one usually

sees many extra lines in the x-ray diffraction patterns, but

80
Table 10

X-ray Diffraction Pattern of Q Phase V Pt

Orthorhombic AuCd-type Structure

ao

= 4 - 413 A, bQ = 2.693 A, c Q = 4.767 A

d(obs)

(a)

d(calc) Relative Intensity

(i) Observed Calculated + hkl

4.741 4.766 w 13 001

3.218 3.238 m+ 61 101

2.674 2.693 mw 22 010

2. 528 vvw

2.370 2.383 mw - 17 002

2.332 2.345 mr 18 Oil

2. 195 2.206 m+ 38 200

2.089 2.097 m 42 102

2.062 2.071 s 100 111

2.002 -- 3 201

1. 777 1. 785 w 7 012

1.699 1. 707 m+ 12 210

1. 650 1. 655 m 12 112

1.611 1. 619 mw 10 202

1.600 1. 607 mw 11 211

1. 582 1. 589 vw 5 003

1. 495 -- 1 103

1. 401 1. 405 mw 6 301

1. 384 1. 388 mw" 6 212

1. 364 1. 368 m 16 013

1. 342 1. 347 mw 8 020

* 1. 307 -- 1 113

1. 296 -- 1 021

1.287 1. 289 w 5 203

1. 249 1. 252 w 8 302

81

(Table 10 Continued)
Table 10 (Continued)

1. 244 1. 246 vs 19 213

1. 243 vs 8 121

1. 192 <1 004

1.1669 1. 1723 vw 4 022

1. 1605 1. 1629 s 19 213

1.1504 ms 8 104

1.1471 1. 1494 ms 10 220

1. 1351 ms 4 312

1. 1309 1.1330 ms 12 122

1. 1173 1 221

1. 1013 1. 1032 mw 4 400

1. 0897 -- <1 014

1.0794 -- <1 303

1. 0747 -- <1 401

1. 0566 1. 0579 mw 7 114

1. 0485 __ 1 204

1. 0336 1. 0353 mw 6 222

1 0 0265 1. 0272 vw 3 023

1.0190 1. 0208 w 3 410

1. 0019 -- <1 313

0.9995 1.0011 mw 3 402

1. 0005 mw 1 123

0.9982 mw 3 411

0.9770 <1 214

0.9702 0.9723 mw 5 321

0. 9533 __ 1 005

0.9384 -- 3 412

0.9318 1 105

0.9305 0.9312 mw 5 223

0.9252 0.9259 mw 7 304

0.9158 0.9168 m+ 10 322

0.9058 0.9062 w 3 403

(Table 10 Continued)

82
Table 10 (Continued)

0.8987 4 015

0.8976 0.8977 m“ 2 030

0.8924 1 024

0.8822 2 031

0.8807 0.8806 m 7 115

0.8756 s 7 314

0.8751 s 3 205

0.8744 0.8747 s 15 124

0.8678 -- 3 501

0.8641 0. 8650 m 13 131

0.8588 0.8588 m+ 16 413

0.8530 0.8533 mw 9 420

0. 8422 <1 323

0.8401 1 032

0.8394 0. 8400 vw 2 421

0.8323 m 10 215

0.8319 0.8315 m 4 230

0.8276 w 7 502

0o 8273 2 224

0.8255 0.8260 s 17 511

0. 8252 s 5 132

0. 8191 0.8191 vw 5 231

0.8095 -- 1 404

0.8033 0.8034 mw 10 422

0. 8000 -- 1 305

0.7943 0.7944 vw 6 006

0.7911 0.7911 vw 7 512

0.7851 vw 8 232

0.7819 -- 3 106

0.7814 0.7816 ms 22 033

w 10 025

-- <1 414

__ „ 503

(Table 10 Continued)

83
Table 10 (Continued)

+ Relative intensities are calculated assuming complete ordering

of the vanadium and platinum atoms; two platinum atoms are in

positions 2(e) at V4 0, ze ; 0, l-ze ; and two vanadium atoms

are in positions 2(f) at V4j Va , zf; ^4 , V2j 1-zf. The best intensity

agreement was obtained for ze = 0. 180 ± 0. 005, Zf = 0. 660 ±

0.005 (standard deviations are estimated). Variations in the

degree of ordering and increments of 0.01 in the z parameters

were considered but in each case they produced a less satisfactory

agreement between observed and calculated intensities than

shown above. The observed and calculated intensities are not

corrected for absorption which is appreciable in the sample.

Comparison of relative intensities was therefore primarily

confined to lines whose Bragg angles (or d-values) were not

greatly different.

84
these lines are not seen after an equivalent heat treatment in

our tantalum- strip high vacuum annealing furnace, or when

the alloys are sealed in pure platinum tubes under a partial

pressure of pure argon0 For example, when proper precautions are taken during annealing of the alloy V Pt, we have

observed an x-ray pattern showing only the lines of a AuCd

type structure, but when the same alloy is annealed in a quartz

tube which was either evacuated to 10~® mm Hg or filled with

a partial pressure of pure argon, then one obtains many extra

lines, presumably due to a contamination of the sample by

silicon. Further details are given in Section IV, F-20

10. Intermediate Phase 3

Greenfield and Beck^ were apparently the

first investigators to report the existence of an intermediate

phase having the A15 ( $-W) crystal structure in the vanadiumplatinum alloy system0 They did not attempt to determine its

range of homogeneity, however. We have observed a broad

homogeneity range for this phase extending from 67. 5 to 8L0

atomic percent vanadium ±1.0 atomic percent. The extent of

this composition range was established mainly by electron

probe microanalyses of two-phase alloys annealed at various

temperatures and rapidly cooled. These results were in good

agreement with data obtained by means of x-ray diffraction and

metallographic examinations. The lattice parameters of this

cubic structure varied with composition from aQ

= 4.813 A (Vrich) to aG = 4.831 A (Pt-rich), and it was further established

by density measurements that the large deviations from the

"ideal" (A^B) composition were not due to atom vacancy forma0

Q4')

tion on the lattice sites. v ’ An A15 type phase having a similarly broad composition range has been reported for the vanadium-

(15) iridium alloy system. v ’ An x-ray diffraction pattern for the

3 phase V^Pt is presented in Table 11.

85
Table 11

X-ray Diffraction Pattern of 3 Phase V^Pt

Cubic CrgS i-type Structure

ao

= 4.817 A

d

(obs)

(A)

Relative Intensity

hkl Observed Calculated +

3 0 403 s 67 no

2 e 405 ms 45 200

2 e 152 m 25 210

1 o 965 vs 100 211

1.701 mw 10 220 22

1. 528 m 15 310

-- -- 3 222

1 o 339 w 5 320

1.290 s 51 321

1 o 206 m 10 400

1. 137 m 9 411, 330

1.079 m t 16 420

1 o 052 mw 5 421

lo 028 m t 15 332

0.9848 mw 5 422

0o 9456 m + 18 510, 431

0„ 8954 mw 7 520, 432

0o 8801 ms 31 521

0 o 8521 m 17 440

Oc, 8266 m 15 530,433

0o 8029 ms 33 600, 442

0. 7922 w 5 610

Go 7817 vs 134 611, 532

Relative intensities were calculated assuming complete

ordering of the vanadium and platinum atoms on the available lattice sties; two platinum atoms in positions 2(a)

000, x/

a

X/s

1/

2 and six vanadium atoms in positions 6(c)

(0% V4 Vg)^, 3

/*,

1 ^)^ 0

l /4 ) (V4 V8 0)(% x /2 o). The

86
calculated intensities do include an absorption correction

and the absorption is appreciable in our sample,, The

observed relative intensities are therefore suitable for

comparison only among lines whose Bragg angle (or dvalues) are not greatly different.

87
-t-> ft

TO Cl ft ft ft > >i > > > ft

d GQL v_n oO CD >-

88

Fig.

21

-

Powder

X-ray

diffraction

patterns

of

single

phase

V-Pt

alloys

obtained

with

nickel-filtered

copper

radiation

(Reduced

53%

for

reproduction.).
11 . Eutectoid: y

«=* 0 + 6

The existence of this eutectoid transformation

was established by x-ray diffraction studies on alloy powders

annealed in evacuated quartz tubes and quenched from various

temperatures. The eutectoid transformation temperature was

located at 970° ± 10°C, and the eutectoid composition at 28» 5

± 1 atomic percent vanadium.

12. Eutectoid: y

<=* 6 + C,

This eutectoid transformation was established

by both x-ray diffraction and metallographic studies on alloys

quenched from various temperatures. A typical eutectoid

microstructure is shown in Fig 0 18a. The eutectoid temperature was found to be 980' ± 10° C, and the eutectoid composition

was established at 39. 0 ± 1 atomic percent vanadium.

13. Eutectoid: y

<=> g + [3

The existence of this eutectoid reaction was

established entirely by metallographic studies on alloys which

were annealed at various temperatures followed by rapid cooling in our tantalum- strip resistance furnace. The alloys were

cooled by turning off the furnace power. The alloy VggPt^g

was particularly helpful for establishing the eutectoid reaction

temperature, since this alloy was composed entirely of a finely

decomposed structure having a feathery appearance when cooled

from above the eutectoid temperature as shown in Fig. 22b . When this alloy was annealed below the eutectoid reaction temperature, however, and rapidly cooled, we observed the spheroidized

structure shown in Fig. 23a. in this manner, it was possible to

bracket the eutectoid temperature by a series of annealing treatments and the eutectoid reaction temperature was thereby located

at 1410 ± 10°C. The eutectoid composition was located by means

of metallographic observations at 55.0 ± 1 atomic percent vanadium.

89
130X (a)

53 at. % V, 47 at. % Pt equilibrated at 1430°C. Alternating bands of C phase (light) and former y phase

eutectoidally decomposed (dark).

80X (b)

55 at. % V, 45 at. % Pt equilibrated at 1430° C and

cooled rapidly. Dark feathery structure is eutectoidally

decomposed y phase.

The light area is a small amount of retained

primary y phase.

Fig. 22 - Microstructures of vanadium -platinum alloys.

90
208X (a)

55 at. % V, 45 at. % Pt equilibrated at 1400° C.

Spheroidized particles of f3 phase in a matrix of

C phase.

SOX (b)

57 at. % V, 43 at. % Pt equilibrated at 1600° C.

Particles of 0 phase in a matrix of prior y phase

which has undergone a eutectoid decomposition

during cooling.

Fig. 23 - Microstructures of vanadium -platinum alloys.

91
1000X (a)

57 at. % V, 43 at. % Pt equilibrated at 1500° C. Particles of 3 phase in a matrix of eutectoidally decomposed prior y phase.

2Q8X (b)

57 at. % V, 43 at. % Pt equilibrated at 1100° C. Dark

areas are 3 phase. Light areas are C phase.

Fig. 24 - Microstructures of vanadium -platinum alloys.

92
14. Eutectic: L Y + 3

This eutectic reaction was located at 1720

± 10° C by melting point measurements.

The eutectic composition was estimated at

63. Oil atomic percent vanadium by examining the microstructure of specimens which had been used in the melting point

determinations. A typical eutectic structure is shown in Fig.

25a.

15. Peritectic: L + « <=> [3

The existence of this peritectic reaction was

established mainly by melting point determinations. Alloys

containing 70 to 90 atomic percent vanadium melted rather

sluggishly, and it was possible to determine the temperature

at which melting was first observed and also the temperature

at which melting was essentially completed. All melting point

measurements on these vanadium-rich alloys were conducted

using high purity thoria crucibles. Metallographic evidence

for the peritectic composition was difficult to obtain, but it

appears that it is located at 81. 5 ± 1 atomic percent vanadium.

93
400X (a)

59 at. % V, 41 at. % Pt "as-cast” condition. Primary

y grains have partially transformed via a eutectoid reaction. Eutectic structure surrounding the primary

grains is y + p.

160X (b)

59 at. % V, 41 at. % Pt equilibrated at 1300° C. Dark

areas are 3 phase. Light areas are C phase.

Fig. 25 - Microstructures of vanadium -platinum alloys.

94
150X (a)

63 at. % V, 37 at. % Pt equilibrated at 1500°C. Islands of eutectoidally decomposed y phase (dark) in

a matrix of p phase (light).

•*. .V.'* -v * ° o • * V .

• c ( f; ... * ej .•> * <=» • A'^r *

. . -

-

' '

160X (b) 83 at. % V, 17 at. % Pt equilibrated at 1030° C.

Spheroidal particles of aV in a matrix of 0 phase.

Fig. 26 - Microstructures of vanadium -platinum alloys.

95
BOX (a)

87 at. % V, 13 at. % Pt equilibrated at 1300° C. The

two phases are p phase and aV.

160X (b)

87 at. % V, 13 at. % Pt equilibrated at 1100° C. The

two phases are p phase and aV.

Fig. 27 - Microstructures of vanadium -platinum alloys.

96
4OX (a)

92 at. % V, 8 at. % Pt equilibrated at 1100° C. Large

grains of aV with a small amount of p phase precipitated at grain boundaries.

Fig. 28 Microstructures of vanadium-platinum alloys.
REFERENCES

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

Binary Systems of Certain Transition Elements, Trans 0 AIME, 206, pp. 265-276, (1956).

2 C Schubert, K. , Frank, K. , Gohle, R., Maldonado, A.,

Meissner, H, G. , Raman, A. and Rossteutscher, W.

Some Structure Bata for Metallic Phases,

Naturwis sense haften, 50, p 0 41, (1963).

3. Maldonado, A. and Schubert, K. , Structure Investigations

in some of the Homologous and Quashiho mologons

Alloy Systems, Z e Met., 55, pp. 619-626, (1964).

4„ Waterstrat, R. M., Crystal Structure Transformations in

VPt Produced by Plastic Deformation at Room Temperature,

Trans , Met. Soc , AIME , 245, pp. 1360-1361, (1969).

5. Sullivan, T„ A., Electrorefining Vanadium, J. of Metals , 17, pp. 45-48, (1965).

6. Dwight, A. E,, Downey, J„ W, and Conner, R. A,, Some

AB3 Compounds of the Transition Metals, Acta Cryst, , 14,

pp, 75-76, (1961).

7. Beattie, H. J,, Close- Packed Structures, in: " Inter metallic

Compounds", pp, 144-165, John Wiley and Sons, Inc., New

York, (1967).

8 . Peters, E. T,, Ryan, J. J., Vilks, I. and Tanner, L. E.,

Mechanical Behavior of Intermetallic Compounds, ASD TDR,

Part m, pp. 62-1087, (1964).

9. Tanner, L. E. , The Ordering of Ni^V, Phys. Stat. Sol., 30,

pp. 685-701, (1968).

10. Giessen, B. C. and Grant, N. J. , The Crystal Structures

of VN12, VPd2, VPt2 and Related AB2 Phases, J. Less

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

98
11. Valentich, J., Thermal Expansion, Measurement Instruments

& Control Systems, 42 (Oct), pp. 91-94, (1969).

12. Steeple, H. , The Crystal Structure of the CadmiumMagnesium Alloy, CdMg, Acta Cryst. , 5, pp 0 247-249,

(1952).

13. Olander, A., The Crystal Structure of AuCd, Z. Krist.

83, pp. 145-148, (1932).

14. Waterstrat, R. M0 and van Reuth, E. C., Effects of

Compositional Variations on the Atomic Ordering in A15

Phases, Proc . Third Bolton Landing Conf. on Ordered

Alloys , pp. 95-110, published by Claitor’s Publishing

Division, Baton Rouge, La., (1970) o

15. Giessen, B. C., Dangel, P0 Nc and Grant, N0 J., New

Phases in the Vanadium-Iridium System and a Tentative

Constitution Diagram, J. Less Common Metals, 13,

pp. 62-70, (1967).

99
V. CHROMIUM-IRIDIUM CONSTITUTION DIAGRAM

A. Previous Studies

In 1955, Raub and Mahler

^

published the results of

their investigation of ten different alloys spanning the entire composition range in the chromium- iridium alloy system. The existence

of two intermediate phases was established. One of these phases

was identified as having a close-packed hexagonal crystal structure

and was stable over a wide composition range. The other intermediate phase was identified as possessing an A15 ( 3 -W) type

crystal structure and was stable over a more restricted range of

composition. A subsequent study of five different chromiumiridium alloys by Knapton^ essentially confirmed the findings of

Raub and Mahler, but the observed lattice parameters were not

in good agreement with those of the previous study. Knapton also

reported the existence of an unidentified phase in an alloy containing

25 atomic percent iridium after it had been annealed at 1200° C.

The information obtained in these two investigations

was very limited in nature, and was therefore insufficient to

permit the establishment of a constitution diagram. No further

work on the chromium-iridium system has since been reported.

B. Materials

The chromium metal used in this study had a nominal

purity of 99. 999% with respect to nongaseous impurities, but it contained approximately two atomic percent oxygen. The presence

'

of oxygen in our experimental alloys was manifested by the appearance

of oxide inclusions (presumably O2O3 ), and it seems probable that

most of the oxygen in these alloys exists in this form. It would be

possible to remove most of the oxygen from the alloys by heattreating either the alloys or the chromium metal used to prepare

the alloys in an atmosphere of dry hydrogen. However, no deoxidation

treatment was employed in this study since it seemed unlikely that

the presence of this amount of oxygen would shift the location of the

100
binary phase boundaries to any significant extent. This conclusion

was supported by our previous experience in determining phase

boundaries in the Cr-Pt system.

The iridium metal used was in the form of powder

having a nominal purity of 99.95%. This powder was compressed

into 1/2-inch diameter pellets before melting them together with

the chromium metal droplets. Spectrographic and chemical analyses

of both of our starting materials are listed in Table 12.

C. Alloy Preparation

All experimental alloys were prepared by arc- melting

the appropriate quantities of each pure element under an inert

atmosphere of 50% argon and 50% helium. Considerable weight

losses occurred during melting, particularly for alloys having

high iridium contents. This is presumably due to the high vapor

pressure of chromium at the temperatures needed to melt these

alloys. Losses were minimized to some extent by placing the

iridium pellets on top of the chromium droplets so that the iridium

would receive the initial impact of the electric arc. Nevertheless,

metal losses during melting were usually about 2% and occasionally

as high as 4%.

An alloy having a nominal starting composition of

Crsoirgo was submitted to a chemical analysis after melting in

order to determine whether the weight losses were caused by

vaporization of both elements or primarily by the loss of only one

component. This alloy had a weight loss of 2 0 6% during melting,

and the chemical analysis revealed that this entire loss could be

attributed to vaporization of Cr. All of our alloy compositions

were therefore determined by measuring the total weight loss

during melting and attributing the loss entirely to a depletion of

the chromium content.

Each sample was melted at least four times and was

inverted in the hearth after each melt. The melting chamber was

101
Table 12

Chemical and Spectrographic Analyses of the

Material Used in Preparing the Experimental Alloys +

Chromium Iridium

Impurity Leico Metals, Inc. Engelhard Industries, Inc.

Ca <10 <1

Fe <10 <1

Mg <10 <1

Ni <10 ND

Pb <10 ND

Pt ND <10

Rh ND <10

Si <10

$

<10 H*C 4 16

* !

295

0* 6540

N* 6

t Values are given in ppm.

ND Not detected spectrographically

* Values obtained by vacuum fusion analysis

No other elements were detected spectrographically.

102
evacuated to a pressure of 10“ ^ mm or less between each melting,

and was flushed several times with the high purity argon- helium

gas mixture. Special precautions were taken with brittle alloys

to insure that no fragments of the alloy were lost if the sample

shattered under the arc.

D. Composition Determination of the Melts

Alloy compositions were determined by chemical

analysis, electron probe microanalysis, or by monitoring weight

losses during melting. The chemical analysis of our alloy C^qI^q

had established that weight losses during melting for this alloy could

be attributed entirely to the vaporization of chromium0 We have

therefore calculated all of our alloy compositions by assuming that

the observed weight losses were due entirely to a loss of chromium.

This procedure yielded alloy compositions which were quite

compatible with compositions determined by electron probe microanalyses. Calibration of the electron probe was accomplished by

using samples of pure chromium and pure iridium together with a portion of our sample CrsQlrgQ which had been chemically analyzed.

Compositions are therefore considered to be reliable within an

experimental uncertainty of ± 1 atomic percent.

E. Temperature Measurement

A platinum versus platinum-10% rhodium thermocouple

was used to measure temperatures up to 1500° C. Temperatures

above 1500° C were measured with a tungsten versus tungsten- 26%

rhenium thermocouple. These thermocouples were periodically

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

standard thermocouple had been calibrated at the National Bureau of

Standards. In addition, calibration against secondary standards was

accomplished using the melting points of pure nickel, platinum,

rhodium, and iridium as described in Section IIX-G. A calibration

of temperature versus furnace current was also helpful in establishing

temperatures within the hot zone of the furnace during the various

annealing treatments.

103
F. Thermal Treatments

1. Homogenization

All samples used in this study were initially

homogenized by heating to within 300° C of their solidus temperatures

for several hours. The homogenization treatment was conducted

with all the precautions normally used in the equilibration treatments

so that the samples could be considered as being equilibrated at

the homogenization temperature. Equilibrations at lower temperatures were conducted consecutively in order of decreasing temperature.

The samples were contained in tantalum buckets

suspended by tantalum wire within the tantalum- strip resistance

furnace during the annealing treatments. For annealing treatments

above 1500° C, the tantalum buckets were lined with tungsten strips

to prevent any reaction between the samples and their containers.

This was necessary since Ir and Ta form a low melting eutectic.

The furnace operated at pressures between 5 x 10 “6 and 5 x 10“

mm Hg. There was no significant visible contamination of the

samples by tantalum and this was later verified by electron-probe

microanalysis.

All samples appeared to be well homogenized.

This was checked by electron probe and metallographic examinations

together with observations of the sharpness of x-ray powder patterns

obtained from filings which had received brief ’'strain- anneal” heattreatments.

It was at first anticipated that alloy samples containing high contents of iridium might undergo serious depletion of

their chromium contents when annealed in a high vacuum at temperatures as high as 2000° C. These temperatures were necessary both

from the standpoint of adequate homogenization, and also to establish

the phase relationships at these high temperatures.

Chromium losses at these temperatures, however,

were not excessive during the relatively brief periods of time needed

104
for equilibration and, in addition, it was found that such losses

were confined mainly to a thin layer at the surfaces of the samples

„

This layer was either removed or avoided during subsequent

studies.

2. Equilibration

Samples to be equilibrated at temperatures up

to 1100° C were sealed in quartz tubes (previously baked out and

evacuated) under a partial pressure of pure argon. These tubes

were placed in a platinum- wound furnace capable of being controlled

within a temperature range of ± 1°C. Alloys to be annealed at

1350°C were wrapped in tantalum foil and placed in platinum tubes.

The platinum tubes were then evacuated and filled with a partial

pressure of argon sufficient to keep the tubes from collapsing at

the annealing temperature. The platinum tubes were then sealed

by crimping the end of each tube in a special device capable of

exerting sufficient pressure to cold weld the end of the tube. This

method permitted annealing the samples for prolonged periods of

time at 1350°C in a platinum- wound furnace in air. Equilibration

at temperatures above 1350° C was accomplished using a tantalumstrip resistance furnace operating at a pressure of 5 x 10“ ^ to

5 x 10-7 mm Kg. The samples were suspended in tantalum buckets

but electron probe analysis indicated no contamination of the samples

by tantalum. Temperatures in this furnace were controlled usually

to within ± 10°C, but more accurate control is possible over shorter

periods of time.

Samples annealed in the tantalum- strip furnace

were cooled rapidly from the annealing temperature by turning off

the power to the heating elements. Samples annealed in quartz

tubes were quenched in cold water. The cooling rates in all cases

were apparently sufficient to permit retention of the phases existing

at the annealing temperature. However, atomic ordering in the Ir

solid solution could not be suppressed by rapid cooling and it was

not possible to retain the disordered face-centered cubic structure

for alloys containing between 23 and 32 atomic percent Cr.

105
A summary of the equilibration treatments used

for all alloy samples is given in Table 121.

G. Methods for Determination of Phase Boundaries

1. Solidus Temperatures

The experimental difficulties associated with

solidus determinations in the chromium- iridium alloy system were

formidable. The alloys melt at temperatures up to about 2400°

and at such high temperatures the chromium vaporizes readily.

Chromium losses during a solidus determination may therefore

be quite large and would introduce serious errors in composition

determination, particularly when conventional techniques such as

thermal analysis are employed. Furthermore, the alloys and their

vapors will react with many refractory materials. Thermocouples

are rapidly contaminated by chromium vapor even when they are

enclosed in protective refractory tubes.

In order to prevent excessive losses of chromium,

it is essential that the method used for solidus determination

should minimize the length of time during which the alloys are

exposed to high temperatures, particularly when they are partially

or completely molten. Conventional thermal analysis techniques

are therefore ruled out since such techniques require that alloys

in the molten condition be heated or cooled slowly in order to obtain

reliable thermal arrest data. The use of microscopic methods to

detect incipient melting, however, are attractive since it is required

only that the sample be heated long enough to obtain a uniform

temperature, and this temperature need not produce complete

melting. The amount of molten material present in the sample

can be only a few percent and will still be detectable by this method.

We have, therefore, adopted a relatively crude

but effective method for solidus determination which involves heating

the sample to temperatures in the vicinity of the anticipated solidus

temperature until evidence of partial melting is seen in the alloy

microstructure after rapid cooling. The sample is exposed to the
Table 13

Summary of Equilibration Treatments

for Chromium - Iridium Alloys

Temperatures

(°C)

Time Alloys

(atomic % Cr)

2160° 1/2 hr. 16

2120° 1/2 hr Q 23

2020° 1/2 hr. 36

2000° 24 hrs. 4, 16,23, 28

2000° 2 hrs. 29, 31, 33 0 5

1800° 10 hrs. 16, 23, 28, 29, 31, 33. 5

t—* 00 oo

o 2 hrs. 36,48. 5, 58

1700° 1 hr. 70

1600° 24 hrs. 4, 16, 23, 28, 29, 31, 33 0 5, 36,

48.4, 58,63

1550° 2 hrs. 70, 78,85,90,95

1400° 24 hrs 70, 72, 74, 76, 78,83,85

1400° 13 hrs. 90, 95

1350° 2 weeks 4,16,23,28,29, 31, 33.5,36.

48/4, 58, 63

1100° 1 month 4, 16, 23, 28, 29, 31, 33. 5. 36,

48.4,58,63,70,78,85, 90,95

107
desired temperature for only about one minute and usually only a

small fraction of the sample is ultimately melted. The method

has been described in detail in Section III, G-l.

The refractory crucibles used to support the

sample during a solidus determination were made of pure thoria0 This material is stable at the high temperatures employed but

there was some evidence that a slight reaction had occurred between

the sample and the thoria crucibles. A violet color was noted at

the surface of the thoria where it had been in contact with the alloy

sample, but very little erosion of the crucible had occurred, and it

is felt that contamination of the sample was insignificant. Metallographic studies showed no difference in structure between the surface and the interior of the alloy samples. The nature of the

reaction was not identified.

Contamination of the tungsten versus tungsten26% rhenium thermocouple was apparently not very significant

since only a slight change in emf was noted after the sample had

been exposed at a given temperature during the solidus determinations.

This is noteworthy since chromium vapor readily penetrates most

refractory tubes used in protecting the thermocouple. Our thermocouples, however, were protected not only by a refractory tube of

BeO, but also by an outer sheath of pure tantalum. This outer

sheath is apparently less permeable by chromium vapor than the

ceramic refractories. Although the tantalum may react to some

extent with chromium vapor (no visible reaction was detected), it

appears to be an effective barrier which acts to prevent contamination

of the thermocouple. The tip of the thermocouple is not exposed to

the direct impingement of chromium vapor since it is inserted in a

small hole drilled into the tantalum holder. (See Fig. 2, Section III.

)

The use of this sheathed thermocouple for temperature measurement

was apparently rather successful and suggests further application

of such methods in cases where metallic vapors are present at high

temperatures. Our temperature measurements were also checked

by simultaneous monitoring of furnace current. The furnace current

108
had previously been calibrated against temperature using a standard

thermocouple up to 1500° C and secondary melting point standards

from 150G°C to 2447°C (melting point of pure Ir). The temperature

indicated by the tungsten versus tungsten-26% rhenium thermocouples

was usually in good agreement with temperatures deduced by

reference to a plot of furnace current versus temperature provided

that care was taken to insure that furnace operating conditions

remained constant. In this connection, it is important to maintain

a uniform temperature and flow of water coolant to the furnace.

It is also important to reproduce the geometrical arrangement of

the holder and thermocouple for each run as well as their exact

location within the hot zone of the furnace so that the experimental

conditions correspond closely to those which existed during the

calibration.

2. Liquidus Temperatures

We were unable to devise a suitable method for

determining the liquidus temperatures of the Cr-Ir alloys and

consequently the locations of the liquidus lines are simply estimated

from the general form of the solidus and the location of the invariant

reactions. The results of our metallographic examination of the

solidus samples was also taken into consideration in estimating the

location of the liquidus through application of the lever rule.

3 0 Other Phase Boundaries

The boundaries of intermediate phases were

determined by metallographic studies and by electron probe microanalyses whenever it was possible to use this method. The

boundaries of the two-phase region y + € could not be determined

by either of these two methods, however, and it was therefore

necessary to rely on x-ray diffraction techniques for such alloys.

Alloy powders were obtained from the bulk samples by grinding

them with a dental diamond wheel and sieving out the course diamond

particles from the residue. The alloy powders so obtained were given

109
a strain-relief anneal for a few minutes in a high vacuum at the

same temperature used for annealing the original bulk sample.

Such procedures were apparently quite satisfactory since the

x-ray patterns so obtained contained sharp lines with well resolved

Koc doublets. Variations of the lattice parameters as a function

of alloy composition were relatively small and it was therefore not

possible to determine phase boundaries by this method.

A sufficient number of alloys were prepared in or

near the two- phase y + € region to insure that the phase boundaries

were reliably established. The compositions of the alloys were

determined mainly by attributing the weight losses during melting

entirely to depletion of the Cr contents of the alloys. This method

appeared to be quite reliable. In some cases, the alloy compositions

determined by this method were verified by electron probe microanalysis of certain single phase alloys. (See Table 14.)

Table 14

Comparison of Alloy Compositions Determined by

Weight Loss and by Electron Probe Microanalysis

Nominal Starting

Composition

Atomic % Cr by

Weight Loss

Atomic % Cr by

Microprobe Analysis

CrsoXrvo 28.0 At. % Cr 28.3 At. % Cr

Cr5oIr 50 48.0 At. % Cr 48.4 At. % Cr

Cr0oIr 4O 57 0 7 At. % Cr 58. 1 At. % Cr

no
H. Experimental Results

1. General

The proposed constitution diagram for the chromiumiridium system is shown in Fig0 29. All alloys were examined either

metallographically or by x-ray diffraction methods or both. The

data obtained by these methods was in good agreement with the

results of the electron microprobe analyses.

2. Solidus and Liquidus

For both chromium-rich and iridium-rich alloys, it was usually possible to detect the first signs of melting by direct

visual observations particularly on sharp edges or corners of the

sample. In some cases, melting was indicated by the appearance

of small molten beads which formed on the surface of the sample.

Whenever the first signs of melting were detected the furnace

power was immediately turned off. This permitted the sample to

cool rapidly. The sample was then removed from the furnace and

subjected to metallographic examination. By observing the relative

amounts of unmelted solid and chilled liquid it was possible to

estimate the approximate location of solidus and liquidus boundaries.

In the case of the iridium-rich alloys, there were

visible signs of melting when only about 5% of the sample was liquid

as established by subsequent metallographic examination. The chilled

liquid was usually observed at prior grain boundaries and produced

a rounded appearance in each grain as shown in Fig. 3Ga. Losses of

chromium during a solidus determination were minimized by an

initial rapid heating of the sample to within about 100°C of the

expected solidus temperature and then raising the temperature in

10° increments allowing about one minute for thermal equilibration

following each increment. Thermocouple readings established

whether the temperature had in fact reached a nearly stable leveL

In most cases the entire run did not last more than 15 minutes and

it appears that this may explain why the chromium losses were not

very great.

Ill
WEIGHT

PERCENT

CHROMIUM

3o‘3«fUVy3dW31

112

ATOMIC

PERCENT

CHROMIUM
160X (a)

16 at. % Cr, 84 at. % Ir heated to 2250°. Melting has

occurred at the grain boundaries.

8 OX (b)

28 at. % Cr, 72 at. % Ir equilibrated at 2000°C. Large

grains of y Ir showing twinning. Small dark pits are believed to be artifacts produced by polishing and etching.

Fig. 30 - Microstructures of chromium -iridium alloys.

113
In view of the considerable experimental problems

in conducting these studies it is felt that the solidus determinations

were reasonably accurate but it is difficult to estimate the probable

error. At best the error would be ± 10° C and it seems likely that

the maximum error would be no more than ± 50° C.

3 0 Iridium Terminal Solid Solution -y and y

1

The solubility of chromium in face-centered cubic

iridium extends to 28 atomic percent chromium and is essentially

temperature- independent. It has been previously established that

atomic ordering occurs in the vicinity of 25 atomic percent Cr^

and that the ordered alloys are ferromagnetic.^5 ^ The critical

temperatures of the order-disorder transformation have not yet

been established but they appear to be strongly compositiondependent. It was found that rapid cooling after arc- melting did

not suppress the ordering reaction in alloys containing 23 to 28

atomic percent Cr but that in the alloy containing 16 atomic percent

Cr atomic ordering could be produced only by annealing the alloy

powder for several days at temperatures between 600° and 700° C.

The ordered phase (y

1 ) possesses a CugAuCL^)

type structure and it appears that the order -disorder transition

temperatures could perhaps be obtained by high- temperature x-ray

diffraction studies.

4. Chromium Terminal Solid Solution - «

The maximum solubility of iridium in body-centered

cubic chromium is about 13 atomic percent Ir at the eutectic

temperature (1670°C). The solvus line shows a pronounced

temperature- dependence and there is evidence of precipitation

of a second phase in alloys containing more than 5% Ir even after

rapid cooling. X-ray diffraction patterns obtained from powders

of these alloys show diffuse or broadened lines of the disordered

114
body-centered cubic « solid solution. The line broadening was not

eliminated by a "strain-anneal" type heat treatment.

The presence of the precipitate in our alloys

accounts for a slight discrepancy between the location of the solvus

boundary as determined by metallographic studies and the boundary

determined by electron microprobe analysis. In this case, the

metallographic data is more reliable since the electron beam

measures a small percentage of the precipitated phase and thus

yields composition values for the « solid solution which are slightly

too high in Ir content. In alloys annealed at 1100° C, however, both

the metallographic and electron probe data indicate a solubility

of about 6% Ir in the chromium terminal solid solution.

5. Intermediate Phase €

Raub and Mahler

^

were the first investigators to

report the existence of an intermediate phase having a disordered

hexagonal close-packed structure in the Cr-Ir system. They

reported that this phase was stable over a broad composition

range and the present study essentially confirms these findings.

The iridium-rich boundary of this phase lies at

about 32 atomic percent Cr and is nearly temperature- independent.

It is perhaps noteworthy that alloys in this region are neither ordered

nor ferromagnetic whereas the face-centered cubic y phase alloys at

28 atomic percent Cr are strongly ordered and also ferromagnetic.

Crystal structure therefore seems to play an important role in

permitting the occurrence of ordering and ferromagnetism in the

Cr-Ir alloy system.

The chromium- rich boundary of the € phase extends

to a maximum chromium content of about 68 atomic percent Cr at

1750° C but the solubility is temperature- dependent and decreases to

60 atomic percent Cr at 1100°C. Variations in the lattice parameters

of the € phase as a function of composition are illustrated in Fig. 31.

115
parameter

(A)

Axial

Ratio

c/

4.31

4.29

4.25

Fig. 31 - Lattice parameters of € phase Cr-Ir alloys.

116

-

parameter
6. Intermediate Phase 3

Raub and Mahler^ first reported the existence of

a phase having the A15 ((3-W) crystal structure in the chromiumiridium alloy system and they reported that it exists under a melting

point maximum. We did not observe a melting point maximum. On

the € -phase side we observed a peritectic reaction. A eutectic

reaction was observed only on the « -phase side rather than on both

sides as reported by Raub and Mahler.^ It is perhaps noteworthy

that Raub and Mahler stated that the eutectic structure which they

observed on the € -phase side was ’’less distinct” than the eutectic

structure which they observed on the « -chromium side. Since they

presented no photomicrographs for their Cr-Ir alloys one might

therefore question whether the structure which they had reported

was indeed a eutectic structure.

The composition range of the 3 phase was determined

by electron probe microanalyses of two-phase alloys quenched from

1550°, 1400° and 1100°C. At 1400°C the composition range of the 3

phase extends from 73 atomic percent Cr to 82 atomic percent Cr

and the lattice parameters vary from ao = 4.689 to ao = 4. 659 with

increasing chromium content.' ' For alloys annealed at lower

temperatures Raub and Mahler^ report a composition range for

this phase extending from 72. 5 to 78. 5 atomic percent Cr with a

similar variation in lattice parameters. The two investigations are

therefore essentially in agreement except for the location of the Crrich boundary of the 3 phase and the questionable melting point

maximum for this phase.

We found no evidence of the unidentified phase

observed by Knapton^ in his alloy containing 25 atomic percent

Ir which had been annealed at 1200°C. Our alloy Cr7glr 22 yielded

an x-ray diffraction pattern containing only the lines of the 3 phase

after annealing at 1100° C for one month. Our alloy Cr^olrso

yielded a pattern showing that the major constituent phase is the 3

phase with a relatively minor amount of the € phase. No other lines

were observed for this alloy annealed at 1100° C for one month.

117
80X (a)

29 at. % Cr, 71 at. % Ir equilibrated at 1800° C. Alloy

contains y Ir (twinned) plus a small amount of € phase.

160X (b)

31 at. % Cr, 69 at. % Ir equilibrated at 1800° C. The

columnar grains showing twinning are presumably y

Ir. Untwinned grains are € phase.

Fig. 32 - Microstructures of chromium -iridium alloys.

118
13 OX (b)

63 at. % Cr, 37 at. % Ir equilibrated at 1100° C. The

matrix is € phase. Particles of (3 phase have precipitated at prior grain boundaries and within the grains.

Fig. 33 - Microstructures of chromium -iridium alloys.

SOX (a)

36 at. % Cr, 64 at. % Ir equilibrated at 2000°C.

Large grains of € phase.

119
160X (a)

70 at. % Cr, 30 at. % Ir equilibrated at 1100° C.

The matrix is p phase containing large grains of

€ phase. Small dark spots are etch pits and

oxide inclusions.

100X (b)

85 at. % Cr, 15 at. % Ir equilibrated at 1100° C. Matrix

is p phase containing second phase particles of aCr.

The distribution of aCr suggests that a prior eutectic

structure may have existed.

Dark spots are oxide inclusions.

Fig. 34 - Microstructures of chromium-iridium alloys.

120
200X (a)

85 at. % Cr, 15 at. % Ir equilibrated at 1550° C. Matrix

is aCr in which precipitation of g phase has occurred

during rapid cooling. Large white areas are g phase.

90 at. % Cr, 10 at. % Ir equilibrated at 1400° C.

Large grains of aCr with a grain boundary network of g phase. Dark particles are oxide inclusions.

Fig. 35 - Microstructures of chromium-iridium alloys.

121
160X (a)

90 at % Cr, 10 at. % Ir equilibrated at 1100° C. 0 phase

formed within grains of aCr and as a network at prior

grain boundaries. Dark particles are oxide inclusions.

Fig. 38 - Microstructures of chromium-iridium alloys.

122
7. Peritectic: y + L <==* €

The existence of this peritectic reaction was

established mainly through measurements of solidus temperatures

particularly for the alloy Cr3iXr69 which lies close to the peritectic

horizontal at 2200° ± 20°C. Alloys annealed within the two-phase

(y + €) region below the peritectic horizontal had a twinned appearanc

(Fig. 32) but there was no clear distinction between v and € phase

regions. X-ray diffraction patterns were obtained from powders

of these alloys and were given a "strain- anneal" treatment at the

temperature of the original equilibration. The patterns clearly

revealed sharp, well-resolved lines of both a face-centered cubic

phase and a hexagonal close-packed phase coexisting.

8 0 Peritectic: € + L <=* 0

This peritectic reaction was also established

mainly by solidus measurements. The results from the alloy

Cr7Qlr 3 o were most helpful in establishing the location of the

peritectic horizontal at 1750° ± 10° C.

9. Eutectic: L ^ + p

The location of this eutectic reaction was established

by extrapolation of both solidus and solvus lines in this composition

region. Metallographie examination of the alloy Crggiris in the

"as-cast" condition failed to reveal a typical lammelar eutectic

structure but this may be due to decomposition of the eutectic

structure during cooling. The direction of adjacent solidus boundaries

definitely suggests the existence of a eutectic reaction.

123
REFERENCES

1. Raub, E. and Mahler, W. , The Alloys of Chromium with

Platinum, Iridium, Rhodium and Ruthenium, Z. Met. , 46,

210-215 (1955).

2. Knapton, A. G. , An X-ray Survey of Transition-Metal

Systems for Sigma Phases, J. Inst. Metals, 87, 28-32,

(1958-1959)o

3. Friederich, E., About Ferromagnetic Platinum- Chromium

and Platinum-Iridium Alloys, Z. Tech. Physik, 13, 59,

(1932).

4. Waterstrat, R. M. and van Reuth, E. C., Effects of

Compositional Variations on the Atomic Ordering in A15

Phases, Proc. Third Bolton Lauding Conf. on Ordered

Alloys , Claitor’s Publishing Division, Baton Rouge, La.,

(1970).

124
VI. CHROMIUM- RHODIUM CONSTITUTION DIAGRAM

A0 Previous Studies

There have been only two previous studies devoted to

the constitution of chromium-rhodium alloys. Both of these

studies were very limited. Raub and Mahler^ studied ten alloys

spanning the entire composition range in increments of approximately 10 atomic percent. They mentioned the existence of only

one intermediate phase having a hexagonal close-packed crystal

structure. Greenfield and Beck^ studied six Cr-Rh alloys and

reported the existence of two intermediate phases; one having a

hexagonal close-packed structure and the other having a 3-W

type structure. These two investigations were not in complete

agreement in regard to the location of certain phase boundaries.

The amount of information available from both investigations,

however, was quite insufficient to establish a constitution

diagram.

B. Materials

Both the chromium metal and the rhodium metal used

in this study had nominal purities of 99. 999% with respect to nongaseous impurities. However, the chromium metal was found to

contain about two atomic percent oxygen. Metallographic

examination of an arc- melted sample of this "pure” chromium

revealed the presence of a small amount of oxide inclusions

(presumably Cr 203 ) but the surface of the sample was remarkably

bright and clean with no visible external evidence of the oxygen

contamination. Nevertheless, when chromium-rhodium alloys

were prepared using this chromium there was a visible green

surface film on the top of each alloy sample. The presence of this

green film on the top of the alloy ingots was initially attributed to

the presence of additional oxygen which had been dissolved in the

rhodium. Accordingly, we attempted to deoxidize the rhodium by

annealing the compressed powder pellets at a high temperature in

125
dry hydrogen and then in a high vacuum. Surprisingly, the green

film persisted. It was finally discovered that the formation of

these visible green films on our Cr-Rh alloys could be prevented

by annealing the chromium in dry hydrogen for four hours at

1700° C prior to melting the alloys. Thus, it appears that the

green film was due to oxygen dissolved in the chromium rather

than in the rhodium. We are unable to explain why the film did

not appear on arc- melted samples of the untreated "pure”

chromium which had not been deoxidized. The arc- melting was

done under identical conditions in both cases using the same high

purity inert gas atmosphere.

As a consequence of the above findings it was decided

to deoxidize all of the alloys which we had already prepared and

to prepare all future alloys using chromium which had been deoxidized by annealing in dry hydrogen. Several of our alloys

were therefore deoxidized by heating them to 1400° C for three

days in an atmosphere of dry hydrogen (see Table IS). All subsequent heat- treatments were then carried out in a high vacuum.

The nominal purities of our starting materials are

given in Table 15.

C. Alloy Preparation

All alloys were prepared by arc- melting the appropriate

amounts of each pure metal under an inert atmosphere of high

purity containing 50% argon and 50% helium. The weight losses

during melting were always less than one percent and were usually

about 0. 5%. We have therefore assumed that all of our alloys have

compositions which correspond to their nominal starting compositions

within the experimental error of our measurements (± one atomic

percent). This assumption was supported by quantitative data

obtained by electron microprobe analyses.

Each sample was melted at least four times and was

inverted in the hearth after each melt. The melting chamber *vas

126
Table 15

Chemical and Spectrographic Analyses of the Materials

Used in Preparing the Experimental Alloys +

Chromium Rhodium

Impurity Leico Metals, Inc.

United Mineral and

Chemical Corp.

Ag ND <1

A1 ND 1 Ca <10 <1

Cu ND 1 Fe <10 4

Mg <10 <1

Ni <10 ND

Pb <10 ND

Si <10 2

C 16 - H* 295 - O* 6540 - N* 6 —

+ Values are given in ppm.

ND Not detected spectrographically.

* Values obtained by vacuum fusion analysis.

No other elements were detected spectrographically.

evacuated to a pressure of 10" 3 mm or less between each melting

and was flushed several times with the high purity argon- helium

gas mixture.

D. Composition Determination of the Melts

Alloy compositions were determined mainly by electron

probe microanalyses and this method was also used to establish

127
that our alloy compositions were substantially the same as their

respective nominal starting compositions,, For example, our

alloy which was nominally CrsoRhgo was found to contain 52.1

atomic percent Cr and 47. 1 atomic percent Rh according to

electron probe analyses. In this case pure Cr and pure Rh were

used as standard reference materials for calibration of the microprobe. In all subsequent calibrations of the electron probe we used

the samples of pure chromium, CrsoRhso, and pure rhodium as

standard reference materials and interpolated the data so obtained

in order to establish the compositions of our other chromiumrhodium alloys. The alloys were also checked for the presence of

impurities which might have been picked up during arc- melting

or annealing, particularly such elements as copper, tungsten,

and tantalum* No contaminants were detected. Compositions are

therefore considered to be reliable within an experimental uncertainty of ± one atomic percent.

E. Temperature Measurement

A platinum versus platinum- 10% rhodium thermocouple

was used to measure temperatures up to 1500° C. Temperatures

above 1500° C were measured with a tungsten versus tungsten- 26%

rhenium thermocouple. These thermocouples were periodically

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

standard thermocouple had been calibrated at the National Bureau

of Standards. In addition, calibration against secondary standards

was accomplished using the melting points of pure nickel, platinum,

rhodium, and iridium as described in Section III-G. A calibration

of temperature versus furnace current was also helpful in establishing temperatures within the hot zone of the furnace during the

various thermal treatments.

128
F. Thermal Treatments

1„ Homogenization

Alloys containing up to 35 atomic percent Cr were

homogenized by annealing them at 1600°C for one hour in a high

vacuum. Alloys containing more than 35 atomic percent Cr were

homogenized at 1400°C for three days in an atmosphere of pure

dry hydrogen which had been passed through a palladium alloy

purifier before being introduced into the furnace0 The homogenization treatments were carried

out using 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 treatments were performed in our

tantalum- strip resistance furnace operating at pressures of between

5 x 10-6 and 5 x 10- ? mm Hg. The samples were suspended in

tantalum buckets suspended by tantalum wire in the hot zone of

the furnace. However, for annealing treatments above 1500°

the tantalum buckets were lined with tungsten strips in order to

prevent any reaction between the sample and its container. This

was necessary since Rh and Ta react to form a low-melting eutectic.

Electron probe studies showed that the samples were adequately

homogenized and essentially uncontaminated. This was also verified

by metallographic examination of the alloy samples and by the

sharpness of x-ray powder patterns obtained from alloy filings

which had been given a brief "strain- anneal” heat treatment.

Loss of Cr by evaporation from the alloy samples

was not a serious problem. Losses of weight during annealing

were quite small (less than 0. 5% in most cases) and it appears

that such losses were confined mainly to a thin layer at the surfaces

of the samples. This layer was either removed or avoided in all

subsequent studies.

129
2. Equilibration

Samples to be equilibrated at temperatures up

to 1200° C were sealed in quartz tubes (previously baked out and

evacuated) under a partial pressure of pure argon. These tubes

were suspended with nichrome wire in a vertical platinum wound

furnace capable of being controlled within a temperature range

of ± 1°C. Equilibration at temperatures above 1200° C was done

in a tantalum strip furnace operating at pressures between 5 x 10“6

and 5 x 1 0-7 mm Hg. Temperatures in this furnace were normally

controlled to within ± 10°C but in some cases it appears that

temperature fluctuations were within ± 5°C especially for those

treatments involving relatively short periods of time.

Samples annealed in the tantalum- strip vacuum

furnace were rapidly cooled from the annealing temperature by

turning off the power to the heating elements. Samples annealed

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

were usually sufficient to permit retention of the phases existing

at the annealing temperature. However, alloys containing the

a -chromium solid solution did show evidence of a decomposition

during cooling. These alloys showed metallographic evidence of

a Widmanstatten- type precipitate within the « -chromium phase

regions but prior structures existing at the equilibration temperature appeared to be readily identifiable. A more detailed description of this behavior is given in our discussion of the experimental

results.

A summary of the equilibration treatments used

for ail of the alloy samples is given in Table 16.

G. 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.

130
Table 16

Summary of Equilibration Treatments

for Chromium-Rhodium Alloys

Temperature Alloys

(°C)

Time (atomic % Cr)

1600° 1 hr. 5,10,15, 20,25, 30,35

1550° 1/2 hr. 90,95

1400° 24 hrs. 5,10,15,20

1400° * 3 days 15, 30, 60, 70, 82. 5, 85

1400° 6 hrs. 40, 50, 78,90,95

1300° 4 hrs. 70, 78,820 5

1275° 6 hrs. 78

1250° 6 hrs. 78

o o O

r— CXI

66 hrs. 70, 75, 78, 82.5

1100° 3 wks. 5, 10, 15, 20, 30

1100° 2 wks B 25,35, 40, 50

1100° 10 days 60, 70,75, 78,82. 5,85,90,95

900° 4 months 5, 10,15,20, 30, 40, 50

900° 2 months 60, 70, 75, 78, 82. 5, 85,90, 95

This annealing treatment was conducted in atmosphere of

pure dry hydrogen.

131
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 in Section III, G-l and is a

rather effective technique when working in the vicinity of the

eutectic compositions or when the solidus and liquidus are

separated only by a small temperature interval as is the case

for most of the Cr-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. 37.

Losses of chromium during a solidus determination were not excessive. This may be partly due to the fact that

the solidus temperatures were below the melting point of pure

chromium and the time needed for a determination was rather

short. The samples were supported in high-purity alumina

(aRR grade) crucibles during the solidus determinations and

there was no appreciable reaction between the metal samples and the

crucibles at the temperatures involved. Nevertheless a slight

violet color was noted on the surface of the crucibles particularly

for alloys rich in c hro miu

0 This violet coloration was similar

to that observed in our studies of Cr-Ir alloys. There was very

little erosion of the crucible and apparently no significant effect

on the alloy samples. The nature of this reaction is unknown.

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.

132
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 temperatures which would place them in close proximity to the boundary

being deter mined„ 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 the case of the 3-phase it was possible to

determine the phase boundaries by quantitative electron probe

microanalyses of two-phase alloys (€ + 3 and 3 + «) annealed at

various temperatures. Alloys annealed within the two-phase <E + «

region, however, could not be used for electron probe microanalyses due to the presence of a fine precipitate in both the € phase and the « 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.

The temperature of the peritectoid reaction was

determined by annealing the alloy Cr78Rh22 at temperatures near

the invariant temperature and then examining the microstructure

133
of the rapidly cooled alloy. Alloys quenched from above the

reaction temperature contained a different major phase than alloys

quenched from below the reaction temperature and in this manner

the invariant temperature was bracketed within the discrete

temperature range shown in Fig. 37.

Phase boundaries in the y + € region were

determined by x-ray diffraction studies of powders obtained by

first annealing the bulk samples at a given temperature and then

removing the alloy powders from these bulk samples with a dental

diamond wheel. Diamond powders detached from the wheel during

grinding were removed by sieving the powder through a - 325

mesh screen. The alloy powders were then given a ’’strain-anneal”

heat- treatment in high vacuum at the same temperature as that of

the ’’bulk” anneal. X-ray patterns obtained from these powders

showed sharp lines with well-resolved Ka doublets at high diffraction

angles. It was not feasible to locate phase boundaries by measuring

lattice parameter variations since the dependence of the lattice

parameters on alloy composition was relatively small and therefore

insensitive.

H. Experimental Results

1. General

The proposed constitution diagram for the

chromium- rhodium system is shown in Fig. 37. All alloys were

examined either metallographically or by x-ray diffraction methods

or both. The data obtained by various methods were all in good

agreement.

2. Solidus and Liquidus

It was usually possible to detect the first signs

of melting in these samples by direct visual observation during

slow heating particularly on sharp corners or edges of the sample.

Melting usually occurred rapidly at a well-defined temperature

and when this occurred, the furnace power was turned off and the

134
WEIGHT

PERCENT

CHROMIUM

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LlJ u

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135
sample was allowed to cool rapidly in the furnace. The sample

was then removed for metallographic examination,, In view of the

relatively low temperatures and sharp melting which were observed

it seems probable that the solidus temperatures have been established to an accuracy of not less than ± 20°C and it is probable that the

error did not exceed ± 10° C. The liquidus temperatures are subject

to more uncertainty and thermal analyses measurements would be

desirable to locate the liquidus temperatures with greater precision.

3. Rhodium Terminal Solid Solution - y

The solubility of chromium in face-centered cubic

rhodium extends to a maximum of about 30 atomic percent chromium

at 1700° C but at lower temperatures the solubility decreases rapidly,.

At 900° C we found a solubility of about 15 atomic percent chromium

in face-centered cubic rhodium and this seems to be in disagreement

with the solubility of eight atomic percent Cr which Raub and Mahler^

reported, presumably on the basis of their alloys annealed at 800° C.

We were unable to retain the cubic y phase by rapid cooling of our

alloy Cr25Rh75 from 1600°C and instead obtained an x-ray pattern

of a hexagonal close-packed structure,, However, we obtained

lines of both the face-centered cubic y and the hexagonal closepacked € phases from a sample of CrSoRh7Q similarly annealed at

1600° C and rapidly cooled. It appears therefore that rapid quenching

is needed in order to retain the y phase in alloys containing more

than 20 atomic percent chromium.

The strong temperature dependence of the (y phase

to € phase) transformation in Cr-Rh alloys is in sharp contrast to

the behavior of this transformation in Cr-Xr alloys where it is

essentially temperature- independent. In this respect, the Cr-Rh

alloys resemble the Cr-Co alloys which also show a strong temperature dependence for the (y phase to € phase) transformation.

It is tempting to speculate that pure rhodium may

resemble pure Co in having a low temperature allotropic transformation to a hexagonal close-packed structure. In the case of pure

rhodium, however, it would appear that the allotropic transformation

136
temperature, if it exists, could not be higher than room temperature and it may therefore have escaped detection due to the difficulty

of obtaining equilibrium conditions at these relatively low temperatures o

4. Chromium Terminal Solid Solution - «

The maximum solubility of rhodium in bodycentered cubic chromium is about 20 atomic percent Rh at the

eutectic temperature (147 5° C). However, the solubility is strongly

temperature dependent below the peritectoid temperature (1265C C)

and decreases to less than five atomic percent Rh at 900°C. This

strong temperature dependence of the solubility is probably responsible for the precipitation effects which have been observed in

<x -phase alloys which have been cooled fairly rapidly.

The precipitate particles are needle- like in

appearance and the alloy microstructure resembles a

Widmanstatten- type structure (Fig. 45a). This structure is also

observed in two-phase alloys annealed above the peritectoid

temperature and rapidly cooled. The prior cx- phase regions in

such alloys apparently decompose during cooling (Figs. 41 to

44). The nature of this precipitation reaction is of considerable

interest since it is apparently responsible for a significant increase

in hardness^ and may also improve the strength of these alloys.

We have therefore obtained x-ray diffraction

data from the surface of a sample containing a sizeable amount

of the needle-like precipitate (Fig0 45a). The data so obtained

indicate that the precipitate particles possess a hexagonal closepacked structure. This being the case, it would appear that the

precipitate is metastable since prolonged annealing of this sample

at temperatures from 90Q°C to 1200°C results in the formation of

the cubic p phase and the complete disappearance of the hexagonal

close-packed structure (Fig. 44). Even more remarkable, however,

is our observation of a lamellar structure in this alloy after a

137
prolonged anneal (two months) at 900°C. This lamellar structure

was found to consist of body-centered cubic « Cr and the cubic 3

phase. The lamellar structure appears to be consuming the prior

needle- like structure (Fig, 433).

The discovery of this lamellar structure suggested

the possible existence of a eutectoid transformation in the Cr-Rh

system but a careful study of all our alloys revealed no evidence

for the existence of a eutectoid reaction. We have therefore concluded that the lamellar structure results from a low temperature

(~900°C) decomposition of the prior metastable needle- like regions.

The needle-like regions form within the large primary grains of

cc Cr during fairly rapid cooling from above the eutectoid temperature (1265° C) or from above the solvus temperature of the « Cr

solid solution. It would be of considerable interest to explore the

mechanical properties of the Cr-rich alloys having the needle-like

precipitate as well as similar alloys annealed at various temperatures so as to produce a variety of microstructures.

5. Intermediate Phase €

The existence of an intermediate phase having a

disordered hexagonal close-packed structure in the Cr-Rh system

has been reported in the previous investigations.^’

^ It has also

been reported that chromium-rich alloys having this structure show

ferromagnetic behavior.^’

^

It has been established in the present investigation

that the solubility limits of the € phase region have a pronounced

temperature dependence. The solubility varies from about 20

atomic percent Cr at 900°C to 68 atomic percent Cr at 1475°C„

A plot of the lattice parameters versus composition is shown in

Fig. 38.

6. Intermediate Phase 3

The existence of an intermediate phase CrsRh

having an A15 (3-W) structure was apparently first reported by

138
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139

ATOMIC

PERCENT

CHROMIUM
60 at. % Cr, 40 at. % Rh equilibrated at 900° C. Large

grains of € phase with precipitated particles of 3 phase.

160X (b)

70 at. % Cr, 30 at. % Rh equilibrated at 900° C. The

clear matrix is 3 phase. Darker areas are € phase

containing an oriented precipitate of 3 phase.

Fig. 39 - Microstructures of chromium -rhodium alloys.

140
160X (a)

70 at. % Cr, 30 at. % Rh equilibrated at 1100°C. The

clear areas are 0 phase. The € phase regions contain globular particles of 0 phase.

70 at. % Cr, 30 at. % Rh equilibrated at 1200° C.

Clear 0 phase regions coexisting with grains of

€ phase which contain some finely precipitated

0 phase.

Fig. 40 - Microstructures of chromium -rhodium alloys.

141
160X (a)

70 at. % Cr, 30 at. % Rh equilibrated at 1300° C. The

clear matrix is € phase containing some acicular particles of 0 phase. The dark grains were formerly aCr

but now contain an acicular precipitate.

130X

70 at. % Cr, 30 at. % Rh alloy in the ’’as-cast”

condition. Eutectic structure of € phase and

a phase. The a phase regions contain a darketching fine precipitate.

Fig. 41 - Microstructures of chromium -rhodium alloys.

142
160X (a)

78 at. % Cr, 22 at. % Rh equilibrated at 900°C. The

clear matrix is 3 phase. Darker grains are decomposed aCr.

160X (b)

78 at. % Cr, 22 at. % Rh equilibrated at 1200° C. The

matrix is 3 phase. Darker grains containing an acicular

precipitate are decomposed aCr.

Fig. 42 - Microstructures of chromium -rhodium alloys.

143
270X (b)

82. 5 at. % Cr, 17. 5 at. % Rh equilibrated at 900° C.

The large clear grains are 0 phase in a matrix of

decomposed aCr. A lamellar growth of a + g is

consuming the finely decomposed a Cr regions.

Fig. 43 - Microstructures of chromium-rhodium alloys.

160X (a)

78 at. % Cr, 22 at. % Rh equilibrated at 1275 °C.

Dark areas containing an oriented precipitate are

decomposed aCr. Light grains are € phase.

144
100X (a)

82. 5 at. % Cr, 17. 5 at. % Rh equilibrated at 1100° C.

Two larger prior grains of aCr have decomposed

to form a two-phase structure consisting of a + 0

phases.

160X (b)

82. 5 at. % Cr, 17. 5 at. % Rh equilibrated at 1200° C.

The matrix is partially decomposed aCr containing

an acicular precipitate. The clear grains are p phase.

Fig. 44 - Microstructures of chromium-rhodium alloys.

145
400X (a)

82. 5 at. % Cr, 17. 5 at. % Rh equilibrated at 1300° C.

The large prior grains of aCr have decomposed to

form an acicular structure consisting of aCr and an

oriented precipitate having a metastable hexagonal

close-packed structure .

400X (b)

85 at. % Cr, 15 at. % Rh equilibrated at 900° C.

Areas which were formerly aCr contain a lamellar

structure of a + 3. Clear areas are 3 phase.

Fig. 45 - Microstructures of chromium-rhodium alloys.

146
8 OX (a)

85 at. % Cr, 15 at. % Rh equilibrated at 1100° C.

Large grains of prior aCr containing precipitated

globules of 0 phase.

160X (b)

90 at. % Cr, 10 at. % Rh equilibrated at 900°C.

Partially decomposed grains of aCr containing

precipitated g phase.

Fig. 46 - Microstructures of chromium -rhodium alloys.

147
160X (a)

90 at. % Cr, 10 at. % Rh equilibrated at 1550° C.

Alloy is single -phase aCr. Large grains show

evidence of polyganization. Dark particles are

oxide inclusions.

160X (b)

95 at. % Cr, 5 at. % Rh equilibrated at 900° C.

Large grains of aCr with a small amount of 3

phase at the grain boundaries. Dark particles

are oxide inclusions.

148
f2) Greenfield and Beck, - * It is difficult to understand why Greenfield

and Beck have attributed the discovery of this phase to Raub and

Mahler* since there is no mention of such a phase in the paper by

Raub and Mahler

The 3 phase apparently forms via a peritectoid

reaction at 1265°C 0 It has a composition range which varies from about

77 atomic percent Cr to 78 atomic percent Cr. It has been reported

to possess incomplete long-range atomic orderingv ' and was found

to be a superconductor at temperatures below 0o 072 K.^~^

7 0 Peritectic y + L «=> €

The existence of this peritectic reaction was

established mainly through measurements of solidus temperatures

particularly for the alloy Cr 35Rh65 which lies close to the peritectic

horizontal at 1700 ± 10° C„ Alloys annealed in the two-phase y + €

region usually produced x-ray diffraction patterns in which sharp

lines of both the face-centered cubic and hexagonal close-packed

structure coexisted. However, it was difficult to discern these

two phases by metallographic studies since the microstructure of

these two-phase alloys appeared to contain only a twinned structure

with little or no contrast attributable to the presence of two phases,

8 . Eutectic L <=> €

This reaction was established primarily by the

observed solidus data and these data indicated a eutectic horizontal

at 1475 ± 10°C.

9, Peritectoid € + « «=* 3

The location of the peritectoid horizontal was

established by means of successive annealing treatments for the

A footnote in the paper by Greenfield and Beck states that

Raub and Mahler ". ..identified the same two intermediate phases

as described here.

0 " but a careful study of the paper which they

referenced shows that only one intermediate phase is mentioned.

This was reported as having^hexagonal close-packed structure

and therefore corresponds to the € phase rather than the 3 phase.

149
alloy CrygRh22» The composition of this alloy is such that when

it is annealed and quenched from temperatures above the

peritectoid temperature its structure contains the « phase as a

major constituent with small amounts of the € phase (Fig.43a)„

When this same alloy is annealed and quenched from below the

peritectoid temperature its major constituent is the 3 phase with

small amounts of the « phase (Fig0 42t} 0 Thus the location of the

peritectoid temperature was established by successive heat-treatments

which rapidly converged on the desired peritectoid temperature,.

The use of thermal analysis methods to detect such a transformation

may lead to serious errors due to nonequilibrium cooling effects

and the present method is therefore more reliable since the alloys

are annealed at each temperature for a sufficient time to insure

that equilibrium conditions essentially prevail,,

150
REFERENCES

1. Raub, E. and Mahler, W. , The Alloys of Chromium with

Platinum, Iridium, Rhodium and Ruthenium, Z. Met., 46,

210-215, (1955).

2 0 Greenfield, P. and Beck, P. A0

, Intermediate Phases in

Binary Systems of Certain Transition Elements, Trans AIME,

206, 265-276, (1956).

3. Raub, E c , Metals and Alloys of the Platinum Group,

J. Less Common Metals, 1, 3-18, (19 59) 0

4. Matthias, B. T. , Geballe, T. H., Compton, V. B. , Corenzwit,

E 0 and Hull, G. W. Jr c

, Superconductivity of Chromium

Alloys, Phys 0 Rev . , 128 , 588-590, (1962)

5. van Reuth, E0 C e and Waterstrat, Rc M. , Atomic Ordering

in Binary A15-Type Phases, Acta Cryst. , B24, 186-196, (1968).

6. van Reuth, E. C., Waterstrat, R. M., Blaugher, R. D.

Hein, R. A. and Cox, J. E., Superconductivity and Long Range

Ordering in A15- Type Structures, Proc. 10th Int. Conf. on

Low Temperature Physics (Ed. Maikov, M, P. ), 137, (1967).

7. Hein, R. A., Cox, J. E., Blaugher, R. D. and Waterstrat,

R. M0

, Superconducting Behavior of A15 Compounds, Solid

State Communications, 7, 381-384, (1969).

8. Blaugher, R. D. , Hein, R. A., Cox, J. E. and Waterstrat,

R. M. , Atomic Ordering and Superconductivity in A15 Compounds

J. Low Temperature Physics, 1, 539-561, (1969).

151
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 Kayvel 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.

152
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 microstructures 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 within

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.

153
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 extending

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 r. p. m. 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. The diamond

154
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

powder in our vacuum furnace at a pressure of less than 10-5 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 as shown in Fig. 21.

155
APPENDIX HI: 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. A comparison

of compositions obtained by the microprobe with chemically

analyzed compositions for several single -phase platinum-chromium

alloys is shown on the next page. The alloy samples were

analysed for both elements and the analytical totals varied from

about 99 to 101%.

156
Electron probe analysis Chemical analysis

atomic % Cr atomic % Cr

28.8 29.7

39.0 39.5

49.2 49.7

59. 1 59.4

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

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 micronprobe 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.

157
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.

158
APPENDIX IV; HIGH TEMPERATURE VACUUM FURNACE

The furnace with which most of this work was carried out

contains four strips of tantalum sheet arranged in a square configuration and heated by a high current (up to 500 amps. ) at a

relatively low voltage (up to about 25 volts) in a vacuum chamber

evacuated to a pressure of about 10 -6 mm Hg. A schematic

diagram of this apparatus is shown below; ^

^g

WINDOW SHIELD

The furnace hot zone has a 2” square cross-section and is

about 5” long with a zone of relatively constant temperature (±10°C)

about 1” long in the center of the hot zone. The four tantalum heater

strips are connected at the bottom to a common tantalum connector

strip while at the top each strip is connected to its own power

terminal by means of four separate tungsten arms. This design

permits unrestrained lateral and vertical thermal expansion of the

159
heating elements resulting in better heat distribution and longer

element life as compared to the cylindrical type of heating elements.

When this furnace is used for a solidus determination, the

apparatus of Fig. 2 is suspended within the hot zone and the tungstenrhenium thermocouple extends vertically to a vacuum connection

which replaces the top viewing window.

This furnace may be operatured at temperatures above

2500° C by replacing the tantalum -strip heating element with a

tungsten-mesh heating element having a similar square configuration. Tantalum radiation shields are similarly replaced

with tungsten shields. In this manner the furnace has been

operated up to the melting point of pure tantalum (about 3000° C).

This required a heating element current of about 700 amps at an

applied voltage of about 20 volts in each leg of the heater circuit.

There are two legs in the heater circuit and the power is applied

pairwise to the four strips resulting in a total power input of

about 30 KVA. The furnace is shown in actual operation in the

accompanying photograph.

160
Fig. 49 - High-temperature vacuum furnace in

operation.

161
APPENDIX V: ATOMIC VOLUME RELATIONSHIPS

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 vs. composition have, therefore,

been constructed for the phases occurring in each of the four

binary systems explored in this investigation. These plots are

shown on the following pages. It is noteworthy that volume

expansions occur in the Cr-Pt, Cr-Ir and Cr-Rh systems whereas

in the V -Pt system, one observes a volume contraction. It is

not yet possible to establish the reasons for this behavior.
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