NISTIR 88-3856

Effect of Interstitial Elements

on Phase Relationships in the

Titanium-Aluminum System

R. M. Waterstrat

U.S. DEPARTMENT OF COMMERCE

National Institute of Standards and Technology

(Formerly National Bureau of Standards)American Dental Association Health

Foundation

Paffenbarger Research Center

Gaithersburg, MD 20899

May 1988

Issued October 1988

Final Technical Report for Period November 1981 to September 1987.

Contracts N00014-82-C-0061 and N00014-86-K-0470

Approved for Public Release; Distribution Unlimited

Sponsored by:

Office of Naval Research

Arlington, Virginia 22209

NISTIR 88-3856

Effect of Interstitial Elements

on Phase Relationships in the

Titanium-Aluminum System

R. M. Waterstrat

U.S. DEPARTMENT OF COMMERCE

National Institute of Standards and Technology

(Formerly National Bureau of Standards)American Dental Association Health

Foundation

Paffenbarger Research Center

Gaithersburg, MD 20899

May 1988

Issued October 1988

Final Technical Report for Period November 1981 to September 1 987.

Contracts N00014-82-C-0061 and N00014-86-K-0470

National Bureau of Standards became the

National Institute of Standards and Technology

on August 23, 1988, when the Omnibus Trade and

Competitiveness Act was signed. NIST retains

all NBS functions. Its new programs will encourage

improved use of technology by U.S. industry.

Approved for Public Release; Distribution Unlimited

Sponsored by:

Office of Naval Research

Arlington, Virginia 22209

U.S. DEPARTMENT OF COMMERCE

C. William Verity, Secretary

NATIONAL INSTITUTE OF STANDARDS

AND TECHNOLOGY

Ernest Ambler, Director
.

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M. TITLE (Include Securitv Classification ) Effect of Interstitial Elements on Phase Relationships in the Titanium-Aluminum System

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1988/November

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16. SUPPLEMENTARY NOTATION

17. COSATI CODES 18. SUBJECT TERMS (Continue on reverse if necessary and identify by block number)

FIELD GROUP SUB-GROUP Titanium Alloys Titanium-Aluminum Alloys

Deutron Diffraction a 2 Phase, T

3A

Phase Diagrams Interstitial Elements

19. A8STRACT (Continue on reverse if necessary and identify by block number)

An apparatus has been designed and constructed for neutron diffraction

studies on polycrystalline samples in a high vacuum at temperatures up to

approximately 1500 °C or higher. Neutron diffraction data have been

obtained from titanium-aluminum alloys containing 12 to 45 atomic percent

aluminum at temperatures ranging from 900 °C to over 1300 °C. Unusually

high background intensities were observed in single-phase a-Ti alloys

containing 14 and 16 atomic percent aluminum between 900 °C and 1000 °C.

This suggests the existence of atomic short-range order as reported by

previous investigators. Addition of 1 atomic percent oxygen to such

alloys was found to eliminate the short-range ordering behavior and

replace it with an extended a-Ti plus a 2 -Ti 3 Al two-phase region.

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Dr. Bruce A. McDonald

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st
19. The cr-Ti solid solution region of the Ti-Al phase diagram was found

to terminate in a peritectoid transf ormat i on at about 1150 °C, but with

increasing aluminum content, it apparently reappears at higher

temperatures. At 1200°C or higher, a single diffraction line is observed

in alloys containing approximately 40 atomic percent aluminum. The line

was tentatively identified as that of a disordered hexagonal close-packed

phase (i.e., a-Ti) ( and this phase has since been reported by other

investigators. The presence of this new phase, which was not anticipated

in previous studies of the phase diagram, implies that the B-Ti phase may

possess a more limited solubility for aluminum than was previously

reported, and this will require that the existing phase diagram be

modified significantly.

Addition of up to 6 atomic percent oxygen to Ti-Al alloys has been

found to elevate both the a-solvus and a-transus temperatures, in

agreement with some previous studies, and it appears that oxygen

stabilizes both the <*-Ti and a 2 -Ti 3 Al phases.

The use of high-temperature neutron diffraction in combination with

other experimental measurements, such as differential thermal analysis,

seems to promise more reliable Ti-Al phase diagrams.
EFFECT OF INTERSTITIAL ELEMENTS ON PHASE RELATIONSHIPS

IN THE TITANIUM-ALUMINUM SYSTEM

Final Technical Report

ONR Contracts NO 0 0 1 4 - 8 2 - C - 0 0 6 1 and NO 0 0 1 4 - 8 6 - K- 0 47

0

November 1, 1981 - September 30, 1987

Submitted by

American Dental Association Health Foundation

Paffenbarger Research Center

National Institute of Standards and Technology

Gaithersburg, MD 20899

Submitted to

Office of Naval Research

Materials Science Division

Arlington, VA 22209
2

FOREWORD

This Final Technical Report is presented by the American Dental

Association Health Foundation, Paffenbarger Research Center at

the National Bureau of Standards, Gaithersburg, Maryland under

ONR Contracts N00014-82-C-0061 and NO 0 0 1 4 - 8 6 - K - 0 47 0 . The work

was sponsored by the Office of Naval Research, Arlington,

Virginia. Dr. B. A. MacDonald was the Program Manager.

This report covers the entire project period from Nov, 1, 1981 to

Sept. 30, 1987. The project was initiated to explore the use of

neutron diffraction as a method for discovering the influence of

interstitial solutes on phase boundaries in the t itanium- aluminum

system.

Dr, R. M. Waterstrat, the Principal Investigator, was associated

with Dr. E. Prince of the Reactor Division at the National Bureau

of Standards in obtaining all of the neutron diffraction data.

Mr. A. Giuseppetti assisted with sample preparation, x-ray

diffraction and optical metallography. We wish to express our

appreciation to Mr. W. Jones and technical staff of the Vacuum

Furnace Systems Corporation, Souderton, Pennsylvania for

designing and constructing the special vacuum furnace used In

this study. We are also especially grateful to Mr. Bill

Lugmayer, P.E. of Lugmayer Associates Inc., Clinton, Maryland for

expert advice and technical assistance in e 1 e c t r on - b e am welding

and to Mr. F. A. Schmidt at the Materials Preparation Center,

Ames Laboratory, Ames, Iowa for fabricating the critical neutron

"window" from special high purity vanadium metal.
3

CONTENTS

Sect ion Page

FOREWORD 2

I INTRODUCTION 4

II CRITICAL EVALUATION OF PREVIOUS STUDIES 6

III ASSESSMENT OF THE BINARY TI-AL PHASE DIAGRAM 10

IV EXPERIMENTAL PROCEDURES 13

A. Sample Purity 13

B. Sample Preparation 18

C. Preparation for Neutron Diffraction 20

Exp e r iment s

V COMPARATIVE EVALUATION OF X-RAY AND NEUTRON 24

DIFFRACTION METHODS

VI DESIGN AND CONSTRUCTION OF A HIGH - TEMPERATURE 25

NEUTRON DIFFRACTION APPARATUS

VII EXPERIMENTAL RESULTS 31

A. The Binary Titanium- Aluminum System 31

B. Short-range Atomic Ordering in the a - T i Phase 32

C. The Ti-Al Phase Diagram up to 50 at. % A1 35

and above 1200°C

D. The Influence of Oxygen on the Phase 38

Boundaries in the Ti-Al System

VIII CONCLUSIONS 42

IX REFERENCES 47

X LIST OF TABLES 50

XI LIST OF FIGURES 51
4

I. INTRODUCTION

The objective of this research was to determine the effects of

interstitial solute elements, such as oxygen, nitrogen and carbon

on the a/a + o. 2 /a 2 phase boundaries in the TI-A1 system. These

effects must be adequately defined since the interstitial

elements are ubiquitous contaminants in titanium alloys beginning

with the initial starting material (titanium sponge) and

accumulating through the various stages of fabrication of the

alloys. Once present, they cannot be easily removed, and there

is justifiable concern that they may produce adverse effects on

the mechanical behavior of these alloys.

The interstitial elements of this investigation are abundant in

the ambient atmosphere during hot forming or thermal welding and

are difficult to exclude. Some contamination of titanium alloys

is, therefore, virtually unavoidable since hot titanium is

capable of scavenging the interstitial elements in an ultra-high

vacuum (pressures < 10" 6 Torr) or from the highest purity inert

gases (impurity contents < 5 ppm). The only practical solution

to these problems is to set maximum allowable limits for

interstitial contamination and then impose effective controls on

it. This requires a knowledge of its effect on alloy properties

and behavior

.

It is already well established that excessive amounts of oxygen,

for example, can embrittle titanium and its alloys. In pure
5

titanium the embrittlement mechanism probably involves an

interaction between dissolved oxygen atoms and dislocations which

produces a ductile/brittle transition (Ref. 1). In titaniumaluminum alloys, however, oxygen also promotes the formation of

the brittle a 2 phase (Ti 3 Al), an atomically ordered version of

the more ductile a phase structure (see Fig. 1). However, there

is a lack of precise information about the effect of oxygen on

the stability of the a 2 phase. Phase relationships in the Ti-Al0 system, for example, have not been well defined.

a Phase a

2

Phase

Fig. 1. Schematic representation of the (aTi) structure and the

q 2 (Ti 3 Al) structure. Open circles represent Ti atoms, dark

circles represent A1 atoms and shaded (grey) circles represent a statistically random occupancy by both Ti and A1 atoms in the

ratio called for by the phase composition. Thick lines outline

the respective unit cell for each structure. The structure of

the oc 2 phase actually involves slight displacements of the Ti

atoms from these "ideal positions" (ref. 10).
6

II. CRITICAL EVALUATION OF PREVIOUS STUDIES

The influence of interstitial elements on phase relationships in

the Ti-Al system has received only limited prior study, and this

was mostly concerned with the effects of dissolved oxygen. The

earliest studies were directed mainly to the influence of

interstitials on the a/a + 0/0 phase boundaries. Van Thyne and

Kessler (Ref. 2) reported that additions of up to 1% oxygen or

nitrogen raise the temperature of the p/a + ft phase

transformation and widen the a + 0 field. They also reported

that carbon is Increasingly soluble (up to about 1%) In the aTi

terminal solid solution as the aluminum content increases to

about 10%, and this increased solubility occurs at progressively

higher temperatures since the peritectoid reaction (a->a + 0 +

TIC) temperatures are also raised significantly.

Schofield and Bacon (Ref. 3) essentially confirmed the results of

Van Thyne and Kessler (Ref. 2) and extended them to higher

aluminum contents. They also presented ternary isothermal

sections at 9 50 °, 1000 ° and 1 050 °C for the Ti-Al-0 system.

Nikitenko (Ref. 4) made the observation that surface oxidation of

titanium alloys at 350 to 500° promotes embrittlement through a z phase formation. Subsequent studies by other authors pursued

this topic in greater detail.
7

Glazova (Ref. 5) presented isothermal sections for the titaniumalum inum - oxyg en system at 800°C and 1000°C, but it is difficult

to reconcile or compare his results with those of other workers

because he claims to have observed several new phases.

Crossley (Ref. 6) has reported that 0.1% oxygen decreases the

solubility of aluminum in primary aTi by about 0.6% (~ 1.0 At. %)

at 700° and thus promotes the formation of the Ti 3 Al (a 2 ) phase.

However, his annealing method raises questions about

contamination, questions which are not fully discussed in his

paper .

Kornilov e t a 1 . (Ref. 7) studied the joint solubilities of

aluminum and oxygen in a-Ti at 600°C, and their data suggest that

the solubility of aluminum in aTi is less than 2% (4 at. %) at an

oxygen content of 5 at . %

.

Vavilova e t a 1 . (Ref. 8) have determined the solubility of

aluminum in titanium to be 6% (by wt . ) at 800°C and have compared

this with the solubility value of 4.5% A1 in an alloy containing

.35% oxygen at the same temperature. They report that the twophase region (a + a 2 )

>

which they say extends from 6 to 15% A1 in

the binary Ti-Al system, is somewhat smaller in the oxygencontaining alloys and that the homogeneity range of the a 2 phase

(reported as 15 to 22% Al) is also smaller for oxy g en - c on t a in ing
8

alloys. They propose, therefore, that oxygen contamination of

the t i t anium- a luminum alloys is linked with decreases in

duct il ity

.

Lim e t a 1 . (Ref. 9) studied the formation of coherent a 2 phase in

titanium alloys containing 8% A1 at 695°C. They report that

increasing the oxygen content from - 500 to 1000 ppm produces a

marked decrease in ductility which then remains essentially

unchanged at higher oxygen levels. This decrease in ductility is

accompanied by a substantial increase in the volume fraction of

the a 2 phase. They suggest that increasing the oxygen content

raises the temperature of the a/a + a 2 transition for a constant

Al concentration. They also present evidence that oxygen

strongly partitions to the a 2 phase and acts as an a 2 stabilizer.

Obviously, more data is needed at various temperatures in order

to define the influence of oxygen on the a/a + a 2 phase

boundaries more accurately. Nevertheless, there is general

agreement among these studies that oxygen promotes the stability

of the brittle a 2 phase (Ti 3 Al) by shifting the a/a + a 2 (a

solvus) phase boundary to lower aluminum concentrations.

The picture of oxygen as an a 2 stabilizer (Ref. 9) is supported

by crystallographic observations. Gehlen (Ref. 10) has shown

that the crystal structure of a 2 is basically the same as that of

the a phase if one ignores the identities of the atoms (See Fig.
9

1). However, in the a 2 phase, the Ti atoms are slightly

displaced from the sites that they occupy in the aTi structure.

The displacement results in larger octahedral interstices which

are more suitable for oxygen atoms. Gehlen (Ref. 10) also

suggests a tendency toward "covalency" as a stabilizing factor

for the a 2 pha s e .

Kornilov e t a 1 . (Ref. 11) employed neutron diffraction methods to

locate the preferred sites of oxygen atoms in the a phase (a

titanium) structure. They reported that the oxygen atoms enter

octahedral interstices in the a phase structure and, in doing so,

displace titanium atoms from their usual positions. Three

different ordering arrangements exist, depending on the oxygen

content. It may be possible to determine the site preference for

oxygen atoms in the a 2 phase (Ti 3 Al) structure in a similar

study. A detailed comparison of the structural relationships in

these two oxyg en - d is s o lving phases is needed to develop a better

understanding of the oxygen partitioning and the influence of

oxygen on the stability. According to Kornilov e t a 1 . (Ref. 11),

the "bonding" strengths in these phases may vary significantly as

a function of the oxygen content.

The general impression that one obtains after surveying all of

this literature is that the influence of interstitial impurities,

such as oxygen, on pha s e boundaries in the Ti-Al system is

inadequately defined. This should not b e surpr i sing since the
10

binary Ti-Al phase boundaries themselves were not well defined at

the time when these studies were done. Each investigator was

apparently forced to establish his own binary "reference

standards" and then determine "shifts" in the location of these

boundaries which were, presumably, produced by a third element.

Unfortunately, the "reference standards" were never adequately

characterized for essentially the same reasons that apply in the

case of the binary alloys. These are discussed in Section III.

As a result, one can seldom be certain whether the observed

"shifts" in phase boundaries are real or simply due to errors in

establishing the true equilibrium reference state in the binary

system.

It should be obvious, of course, that defining the influence of a

third element is much easier if one first obtains an accurate

binary phase diagram. However, Impatience to obtain "practical"

results can cause one to place excessive reliance on existing

binary phase diagrams without seriously questioning their

validity. The reliability of binary Ti-Al phase diagrams is,

however, a factor of critical importance in several large

industries, and the binary system should be re-evaluated before

beginning any study of third- element effects.

III. ASSESSMENT OF THE BINARY TI-AL PHASE DIAGRAM

There is an extensive literature on the Ti-Al alloy system mainly

because of its practical relevance in almost every aspect of
11

commercial titanium applications. The literature, however, is

often contradictory and somewhat controversial. There are now

some excellent critical reviews (Refs. 12-14) of this system. An

"assessed" Ti-Al phase diagram has been compiled recently by J.

L. Murray (Ref. 14) based on her extensive evaluation of about

100 references dating back as far as 1923. Despite the vast

amount of time and effort which have been expended on Ti-Al phase

relationships over the years and the great practical significance

of this system, it is surprising that the details of the phase

diagram are still not adequately established.

The experimental difficulties that one encounters when attempting

to determine phase equilibria in titanium alloys are, for the

most part, widely recognized. Advancing technology has evolved

methods for dealing with these problems, however, and many

reliable titanium phase diagrams currently exist. But the Ti-Al

system, being more difficult to study, is not one of them. The

reason why the Ti-Al phase equilibria are especially difficult is

that certain Ti-Al h igh - t emp e r a tur e structures cannot be retained

by quenching. When these samples are studied at room temperature, severely decomposed structures are not always easily

interpreted. The samples should be studied at the requisite

h igh - t emp e r a tur e s where critical information is most needed.

Unfortunately, many investigators do not have access to

experimental apparatus which is capable of h igh - t emp e r a t ur

e
12

studies on reactive metals. High- temperature methods can also be

tedious and time-consuming, and this too inhibits their use.

Another reason why Ti-Al phase equilibria are poorly understood

is that most of the previous research effort was narrowly focused

on what was believed to be the region of greatest "practical"

relevance, i.e., temperatures below 1200°C and aluminum contents

of less than 10 at. % where one sees the initial slopes of the

transus and a solvus boundaries. The rest of the diagram has, to

a large extent, been considered to be somewhat irrelevant for

"practical’' purposes and not given much attention. However, this

narr owly - focus e d effort has obviously not produced the desired

result. We still do not possess a thorough understanding of the

a solvus boundary or a complete picture of the aTi solid solution

region. Phase boundaries of the a 2 phase are also rather

uncertain, even though their locations have an important bearing

on how much of the a 2 phase will form when the alloy composition

enters the two-phase (a + a 2 ) region. Hurray's review (Ref. 14)

concludes that more work is needed to define coherent vs.

incoherent phase boundaries and to elucidate the effect of

impurities such as oxygen. Our own work will show that there are

important questions to be answered concerning the existence and

extent of short-range atomic ordering in the aTi solid solution.

It will also show that the solubility of A1 in the /3 phase has

been seriously overestimated and that a h igh - t emp e r a tur e phase
13

region at about 40 at. % A1 has probably been missed completely

in the previous studies.

In any case, it seemed obvious to us that before one can

understand the behavior of a third element, it is necessary to

have a better understanding of the binary Ti-Al system.

Therefore, this is where we directed our initial studies.

IV. EXPERIMENTAL PROCEDURES

A. Sample Purity

Commercial titanium alloys have an interstitial element content

of roughly 1000 ppm. It is important to know whether

contamination at this level is having a significant effect on the

microstructure and mechanical properties of these alloys. It

seems desirable, therefore, that a comparison should be made

between samples containing about 1000 ppm of oxygen, for example,

and samples containing lower oxygen contents (< 200 ppm) as well

as with samples having higher oxygen contents ( - 1 at. % 0).

This means that the oxygen content of the samples must be

carefully controlled at each step in the experimental procedures.

It also means that no other impurities should be introduced and

that the aluminum contents of the samples should be constant and

well established if one is to have a valid comparison.

Therefore, throughout this investigation, we placed a special

emphasis on the prevention of sample contamination since

impurities, once introduced, are virtually impossible to remove.

Of course, when conducting experiments on titanium alloys at high
14

temperatures, some contamination is unavoidable because of the

exceptional reactivity of these alloys. Every effort was made to

limit this contamination, and we have conducted numerous chemical

analyses to determine whether it had attained serious levels as a

result of any of our experimental procedures. We believe that

the effort required to accomplish this is justified by the

confidence which we can place in our final results.

Our first task was to obtain high-purity starting materials

having particularly low levels of interstitial contaminants (<300

ppm). This is not a difficult task for metals such as aluminum,

which is readily available in the requisite state of high purity.

However, the elimination of interstitials from pure titanium is

quite difficult, and there are few reliable sources for such

material. The purest titanium ever produced commercially was

probably the so-called ''iodide" titanium sold by the Foote

Mineral Co. , the New Jersey Zinc Co. and other manufacturers in

the 1 9 50 's. According to a Foote Mineral Co. , "Products List"

dated June 1, 1959, this material typically contained 50 ppm

total interstitials (C, N and 0). However, the production of

this material was apparently discontinued some years ago. We

contacted the Foote Mineral Company and were told that the

company no longer makes or sells this product and would not do so

even on special order. Our inquiries to other American

manufacturers brought forth similar responses as did our

inquiries to European manufacturers. Having concluded that it
15

was not available in Western countries, we turned to the "Eastern

bloc." A letter to the Soviet trade mission brought no response,

but we subsequently learned that the People's Republic of China

has recently been exporting "iodide" titanium to the United

States where it was being sold by an American importer. We,

therefore, purchased two bars of this material and submitted some

of this material to vacuum fusion analysis. Metallographic

studies revealed no evidence of particulate matter either in the

grain boundaries or within the grains, and the bars had an

excellent appearance (Fig. 2). Our vacuum fusion analysis of

this bar revealed a total interstitial content (C, N, and 0) of

150 ppm (see Table 1). The only material now being produced in

the United States having a comparably low interstitial content is

probably the TIMET Corporation's "El 60" product which,

unfortunately, contains about 3% residual salt left over from the

electrolytic refining process. This is enough salt to produce a "sputtering" of the electric arc during "melt-down" operations

and users of the product generally conduct a preliminary melting

in order to eliminate it. It is commonly assumed that, since the

"sputtering" is eliminated by this pre-melting, the residual salt

has also been eliminated by simply being vaporized at these high

temperatures. This assumption, however, may be incorrect. We

deliberately added 0.685 grams of dehydrated salt to 13.079 grams

of pure (salt-free) iodide titanium and then melted this (5% salt

+ 95% Ti) mixture four times while monitoring the weight losses
Table 1 MANUFACTURERS' "TYPICAL" ANALYSES FOR

PPM BY WEIGHT)

.

'HIGH- PURITY TITANIUM (IN

1982

1959

Foote ( 2

}

1982

Chinese

Iodide

1982

MRC (

3

>

Element

Electrolytic

TIMET El 60 (1)

Iodide

Process

Process

(Nominal) (Analyzed)

MARZ

Grade

Si 20 50 40 4

A1 10 50 100 5

Mg 10 30 N.D. <.l

Pb N.A. ( 4 5 N.D. ( 5 ) 10 <.l

Mn 10 30 N.A. 5

Cr 10 20 N.A. 3

Sn 200 10 N.A, <.l

Fe 10 20 100 22

Ni 10 10 N.A. 10

Mo 10 N.D. N.A. <.l

Ca N.A. 30 N.A. 1

Cu 10 3 N.A. 8

V 200 N.A. N.A. 5

w N.A, N.D. N.A. <.l

Zr 10 500 10 3

C 90 10 100 40 35

0 100 20 210 100 140

N 20 20 50 10 5

Cl 800 N.A. 30 3

Stated Purity 99.85+ 99.92 99.90+ 99.97

Prices $88/lb $125/lb $57/lb (3)

Year (1982) (1959) (1982)

Sources El 60 Titanium Metals Corp.

P.O. Box 2128

Henderson, Nev. 89015

Phone: 702-564-2544

Chinese Iodide Continental Metals Corp

820 Second Ave

.

New York, N.Y. 10017

Phone: 212-421-9811

MARZ Materials Research Corp

Orangeburg, N.Y,

Phone: 914-359-4200

( 1

5

Somewhat higher purities are also available on special order.

( 2

}

This material is no longer available.

( 3

}

This is a premium product for which prices are quoted on an individual basis

(4) N.A. - not analyzed.

(5) N.D. - not detected.
17

i^i Mi

,1 I GH .PURI TT T I T A N I UM

U M .* M A 9 E IN CHINA

L ;h < «.W

P if * N . W >

fit -V cNOi

'? » K • S

M l J

Fig. 2. High-purity "iodide" Ti crystal bar imported from the

People's Republic of China; used as the starting material of this

inve s tigation

.

Weight Losses During Arc-Melting

Fig. 3. Weight losses produced by repeated arc-melting for pure

iodide Ti, El 60 Ti and synthetic Ti + 5% dehydrated NaCl

mixture

.
18

after each melt. The total observed weight loss after the

melting was 5.6%, which significantly exceeds the 5.0% salt in

the original mixture (see Fig. 3). It seems obvious, therefore,

that these losses are not due to simple evaporations of salt.

They probably involve the Ti also. In further support of this

conclusion, we melted some pure iodide titanium (salt-free) four

times under the same conditions and observed a total weight loss

of less than 0.1%. We, therefore, concluded that our use of the

"El 60" material would not only impose the added burden of premelting but would introduce uncertainties in regard to the salt

content and thus pose an unnecessary risk to the accuracy of our

data. Furthermore, the cost of the "El 60" product was

significantly higher than that of the Chinese "iodide" titanium.

A summary of the various available source materials is given in

Table 1.

B. Sample Preparation

Before melting our samples, we undertook an evaluation of several

alternate melting procedures to determine which procedure would

best maintain the purity of our starting materials. Trial

samples of our "iodide" titanium were melted (four times each)

using these three separate procedures:

Procedure #1 - Evacuate the melting chamber, using only a

roughing pump, to a pressure of ~ 10' 6 Torr

and then introduce special high-purity argon

gas before arc-melting.
19

Procedure #2 - Evacuate the melting chamber using an oildiffusion pump (backed by the roughing pump)

to a pressure of - 10" 6 Torr and then

introduce the special high purity argon gas

before arc-melting.

Procedure #3 - Evacuate the melting chamber as in Procedure

#2 and backfill with the special high-purity

argon gas. Purify the argon in situ by

electrical heating of an interior 1" x 6"

zirconium strip (0.004” thick) followed by

arc -melting

.

The samples prepared by each of these procedures were then

submitted to a va cuum - f us ion analysis with the results shown in

Table 2.

Table 2

Effects of Various Furnace

Atmospheres on Arc-Melted

Titanium

Pure iodide Ti

Roughing Pump (10*4 mmHg)

Diffusion Pump (10*6 mmHg)

Zr-Gettered Argon

Gas Content (ppm)

0 N C

100 10 40

120 10 40

90 10 80

140 10 1260
20

Procedure #1 is quite satisfactory for our purposes. The use of

the o i 1 - d if f us ion pump in procedure #2 did not provide much

additional lowering of the oxygen content and seems to have

increased the carbon content, perhaps by back- s tr earning of the

oil in the diffusion pump. All subsequent sample preparations

were, therefore, carried out using procedure #1. Each 20 gram

sample was melted four to six times and inverted between melts.

Melting losses were always < 1%. Chemical analysis of samples

containing substantial amounts of A1 , 0 , N or C (see Table 3)

showed compositions corresponding closely to those expected based

on the nominal amount of each element that had been added.

Oxygen was added as Ti0 2 , nitrogen as chemically analyzed TiN and

carbon as TiC. High-purity metallic aluminum was also used.

Samples were homogenized using a procedure developed by Shull e t

a 1 . (Ref. 15) who showed that homogenization in the a phase

region avoids the chemical segregation which would occur if the

sample were cooled from the 0 region. All homogenization

annealing treatments were done in a tungsten resistance furnace

in a vacuum corresponding to pressures of < 10“ 8 Torr . Homogenization times were one week at 1000 °C and two weeks at

900°C. The samples were then furnace cooled to room temperature

before being removed from the furnace.

Preparation for Neutron Diffraction Experiments.
21

Each 20 g button was sectioned with an abrasive wheel using

plentiful water-cooling to prevent overheating. Small (3/32"

dia.) holes were drilled in the ends of each half button using

electrical discharge machining (EDM) techniques, and the two 10

gram sections of each sample were then suspended, one above the

other, in the high- temperature neutron diffraction apparatus with

vanadium wire. The vacuum lines and connections in the apparatus

were all metal with copper gaskets so that no organic material

was exposed to the vacuum. The high vacuum was achieved using a turbo -molecular pump which maintained pressures of at, least 10~ 5

Torr even when the chamber was heated to high temperatures. This

seemed to be adequate since all samples had bright, clean

surfaces when removed from the chamber, even after having been

heated to high temperatures in prolonged (overnight) experiments.

Special interlock systems were incorporated to prevent damage to

the apparatus in the event of loss of vacuum, wa t e r - c oo 1 ing or

electric power. A continual check was provided against possible

short-circuiting by the sample against the heating elements by

connecting a resistance meter between them.

Although some preferred crystalline orientation no doubt existed

in our samples, it was not a serious problem for our experiments.

Because the grain size ranged from one to four mm, about - 100

grains were immersed in the neutron beam. The samples were,

therefore, not rotated during the experiments.
22

The initial experiment on each sample involved heating it in the

high-temperature apparatus to a temperature corresponding to that

of its prior homogenization and holding for about 1 hour to

permit temperatures to equilibrate and to disperse any

nonequilibrium structures which may have formed when the sample

was cooled to room temperature after its homogenization

treatment. A complete pattern was then recorded in the angular

region 2 9 — 14 to 42 degrees. Thus, equilibrium data was

obtained for most samples first at 1000 e G and then, after rehomogenizing, at 900°C. In addition, data was taken on selected

samples by temperature scanning. This involved raising the

temperature in small (5° to 10°G) increments while scanning for

the strongest diffraction lines of existing or anticipated

pha s e s

.

One could thereby observe a phase t rans it ion by

f o 1 lowing the d i s app earance of one line and the appearance o f

another

.

We f ound that about 10 to 15 minutes time wa s

suf f ic lent t o re-establish phase equilibrium in our samples a t

temperatures above 1000°C when the temperature was raised by

about 10°C. Thus, phase transformation temperatures could be

established to about ± 5

0 C if sufficient time was allotted

e following each temperature increment. Phase boundaries

determined in this manner agreed well with those determined using

DTA methods within our experimental accuracy (± 10 ° C) .

Vaporization of aluminum from our samples did not pose a

significant problem at temperatures up to 1300 °C. None of our
23

samples has yet been studied above 1350°C. We estimate that our

apparatus is capable of operation up to at least 1500°C. The

maximum temperature is set by the melting temperature of the

vanadium furnace elements (approximately 1930°C). We have

chemically analyzed several of our samples after prolonged study

in our apparatus at the higher temperatures, and it appears that

there is minimal, and probably insignificant, contamination (see

Table 3 )

.

TABLE 3 ANALYSES OF BINARY TI-AL SAMPLES AFTER COMPLETION OF ALL

EXPERIMENTS INCLUDING HOMOGENIZATION AND HIGH -TEMPERATURE

NEUTRON DIFFRACTION EXPERIMENTS.

At. % A1 At. % A1

Sample # (Intended) (Analyzed) 0 (ppm) N (ppm) C (ppm)

1 14.00 14.05 380 4 158

2 16.00 16.03 296 r 142

3 18.00 . 17.80 212 10 143

4 20.00 19.76 302 4 122

5 22.00 21.87 334 20 278

No a 1 1 emp t was ma d e to eva].uat e the peaik intensities

quant itatively although future s tu dies might include such

r e f in emen t s . This study was inte nded to s imply locate various

pha s e bound aries and identify the ph as e s Obviously, however,

phase trans

:

formation temp<era tur e s can b e located more precisely

and at a s ignif icant ly lower co s t using DTA methods (Refs.

13,15 , 16 )

. The primary advantage o f using h igh- temperature
diffraction methods lies in its identif ica t ion of the hightemperature phases wh ich participate in these transformations and

in its unique sensitivity for detecting ordered structures in

this alloy system.

V. COMPARATIVE EVALUATION OF X-RAY

AND NEUTRON DIFFRACTION METHODS

The basic principles of X-ray and neutron diffraction are quite

similar, and both of these methods can be used to identify the

structure of po ly cry s t a 1 1 ine samples. However, there are some

fundamental differences between the two methods in regard to such

things as sample dimensions, absorption characteristics, atomic

scattering behavior, availability of suitable radiation sources

and, last but not least, there is a significant cost

differential. All of these factors must be considered when one

chooses which of these methods should be used in specific

instances

.

X-ray sources are less expensive than neutron sources and are

readily available in most laboratories, whereas neutron beams

having the desired wavelength and intensity are obtained from

atomic reactors. Once having located a suitable neutron source

(most often from an atomic reactor), it is necessary to obtain

the collaboration of the reactor group's staff who will

coordinate and supervise the experiments in accordance with the

normal operating schedule of the reactor which includes
25

alternating "start ups” and "shut downs" for reactor maintenance

or modification. One must schedule experiments well in advance

so that they do not conflict or interfere with other experiments

being done at the reactor.

Experimental apparatus to be used at the reactor face must be

safe and reliable for continuous operation, unattended, over many

hours. It is also desirable that the apparatus can be quickly

placed in operation at the reactor at a pre - s cheduled time and

removed promptly when the experiment is finished to make way for

others. Reactor time is a valuable commodity, and prolonged

delays in using one's allotted time are not welcome. There is

also a lively competition for "reactor time" which generates

pressure to use it efficiently.

VI. DESIGN AND CONSTRUCTION OF A HIGH - TEMPERATURE

NEUTRON DIFFRACTION APPARATUS

Our choice of neutron diffraction for investigating the

constitution of Ti-Al alloys was predicated on several factors.

First of all, it offers an especially sensitive means for

discriminating between the a-phase terminal solid solution and

its ordered counterpart, the a 2 phase (see Fig. 1 and Fig. 4).
Neutron

Diffraction

Patterns

A

=

1.5426

A

o

o o

o O © © © © o © © in o

f*

r 00 m CM 00

in

(sjunoo) Ajisuajuj

15.00

20.00

25.00

30.00

35.00

40.00

Fig.

4.

Neutron

diffraction

patterns

for

(aTi)

phase

(annealed

Ti

powder)

and

a2

(Ti3

Ai)

phase

(annealed

Ti3

Al).

Peaks

are

identified

by

the

Miller

indices

of

their

respective

unit

cells.

The

patterns

were

obtained

at

room

temperature

using

unrotated

samples.
2 7

Secondly, it permits the use of massive samples 20 g.) with

low surf ace - to -volume ratios and irregular shapes. Sample

contamination is, therefore, minimized when the sample is heated

in a vacuum or inert gas atmosphere. Thirdly, the relatively

large sample size insures that a larger number of grains will

intercept the incident radiation beam than is the case with Xrays . This permits a more random sampling of the crystal

orientations. Finally, neutron diffraction offers a wider

latitude in designing associated high- temperature apparatus. One

can surround the sample with furnace components and vacuum

chambers while still maintaining an adequate beam intensity at

the sample. Furthermore, if these components are fabricated from

vanadium, there will be essentially no diffraction peaks from

them and one will see only the diffraction lines from the sample

itself. Pure vanadium is a natural "null matrix" material for

neutrons which contributes no significant coherent diffraction

peaks and has a minimal neutron absorption coefficient.

We, therefore, set out to design and construct an apparatus

having the unique feature that all portions of the apparatus

which intercept the neutron beam would be constructed of pure

vanadium (see Fig. 5). This includes heating elements, heat

shields and vacuum chamber walls. The heating element and heat

shields are located inside the vacuum chamber, along with the

sample, to protect them from oxidation. The neutron beam is

transmitted efficiently into and out of the vacuum chamber by a
28

Furnace for High Temperature

Neutron Diffraction Studies

Tantalum Wlr® Titanium

Fig. 5. Schematic representation of h igh - t emp e r a tur e neutron

diffraction apparatus operable to about 1500 S G under a high

•vacuum .

vanadium "window" (1/8" thick x 2" wide) incorporated directly in

the wall of the vacuum chamber (Figs. 5 and 6). In order to

insure that the window will maintain a high vacuum ( - IQ” 6 Torr),

it is welded to the chamber wall. This means that the wall

itself must be constructed from a material which has a thermal

expansion coefficient and thermal conduction characteristics

close to those of vanadium. Titanium was chosen for these

reasons. The entire apparatus (except the window, etc.) was,

therefore, constructed of commercial -pur ity titanium using

electron-beam welding. After preliminary experiments indicated

that vanadium and titanium could be welded with this method, the

vanadium window itself was welded successfully to the chamber

wall without cracking. A helium leak test indicated that the

completed chamber was leak-tight at pressures down to 10' 9 Torr.
29

Fig. 6. Photograph of the completely assembled high- temperature

neutron diffraction apparatus ready for attachment to the vacuum

sy s t em

.

The temperature of the sample was regulated by means of a control

thermocouple which was placed at the base of the hot zone so that

it would not be in the path of the neutron beam. The control

thermocouple was then calibrated against a standard thermocouple

that was placed at the exact position to be occupied by the

sample. These thermocouples were also calibrated against the

electrical current in the furnace heating elements. Later

experiments showed that the temperature at the sample position

remained constant (± 1°C) during experiments of several hours
30

duration and that the sensitivity of the temperature controller

permitted us to regulate the temperature in increments as small

as 1°C or less. The actual temperature of the sample, however,

seemed to be 30° to 40° higher than indicated by the calibration,

based on a comparison of observed a/0 transitions with those

reported int he literature. The accuracy could, of course, be

greatly improved if the thermocouple were placed near the hot

sample, but then it would be exposed to int er cont amina t ion with

the sample and might also contribute, undesirably, to the

diffraction pattern. We felt, moreover, that these additional

risks were unnecessary for our studies since we were primarily

interested in the variation of the transition temperatures from

one sample to another rather than in the absolute values

.

Furthermore, the absolute values of many transitions are

obtainable, for a given composition, from previous studies

already published (Ref. 14).

We have not yet determined the maximum temperature at which this

apparatus can be operated effectively. It is certainly somewhere

below the melting point of the vanadium heating elements (~

1900°C), but we have so far operated it only up to about 1350°C.

We observed some heating of the chamber walls, but this can

probably be eliminated by design changes and more effective

wa t er - coo 1 ing . In principle, there is no reason why a similar

apparatus could not be operated at temperatures up to 3000°C, but

then one would have to construct the heating elements and heat
31

shields of tantalum or tungsten with a resultant contribution to

the diffraction pattern. A photograph of the apparatus is shown

in Fig. 6

.

VII. EXPERIMENTAL RESULTS

A. The Binary Titanium- Aluminum System

Neutron diffraction patterns obtained at room temperature for

pure aTi phase and for a 2 phase (Ti 3 Al) are shown in Fig. 4.

Both of these phases have essentially the same hexagonal

structure geometry with similar atom positions (Fig. 1). However,

the atom sites are occupied by titanium and aluminum in a

statistically random manner for the aTi phase, whereas in the a 2 phase the atom distribution is highly ordered. This ordering of

the Ti and A1 atoms doubles the "a" parameter while the " c

parameter remains virtually unchanged (see Fig. 4). The effect of

the atomic ordering is to produce strong "superlattice lines”

(100, 101, and 110) in addition to the so-called "fundamental

lines" (200, 002, 201) which remain essentially unchanged. The

neutron diffraction pattern of the a 2 phase, shown in Fig. 4, did

not change significantly between room temperature and 1000°C, and

this suggests that there are no significant structural changes

occurring in this temperature range. One can, therefore, detect

very small amounts of the a 2 phase (- 2-3%) in an a phase matrix,

but the converse (a in a 2 phase) is rather difficult to detect

due to the overlapping lines.
32

We then proceeded to examine Ti-Al alloys containing 14, 16, 18,

20 and 22 at. % A 1 , first at 1 0 0 0 ° C and then, after further

homogenization, at 900°G. The structures which we observed in

each case (a and a 2 phases) and the relative proportions were

essentially in agreement with those indicated in the recent phase

diagram by Shull e t a 1 . (Ref. 13) as well as with the assessed

Ti-Al phase diagram by Murray (Ref. 14). Our data points are

superimposed on the ne wly - pr opo s e d phase diagram in Fig. 7.

B. Short-range Atomic - Or dering in the aTi Phase

It is noteworthy that the samples containing 18 and 20 at. % A1

contained major quantities of the a 2 phase at 900°G. These

essentially disappeared when the sample was heated at 1000°C.

The sample containing 18 at. % A1 produced only the lines of the

aTi phase, but all of the observed lines were immersed in an

unusually high background intensity. A similar high background

was also observed in the aTi pattern of the 16 At. % A1 sample at

both 1000°C and 900°G, but aTi patterns from other samples had a

normal background. Previous workers (Refs. 17 , 18 ) have reported

evidence of short-range atomic ordering in this region of the

phase diagram, but the results were somewhat ambiguous.

We believe that the high background intensities are direct

evidence of short-range atomic ordering. Neutron diffraction

methods should be particularly sensitive to short-range ordering

In this case because the atomic scattering factors of Ti and
33

Oo ejn;BJediu0i
34

A1 atoms are of nearly equal magnitude but opposite sign (b =-

0.38 and + 0.35 respectively). This combination would be

expected to produce a large apparent "temperature factor"in

neutron diffraction data, i.e., high diffuse background

int ens ity

.

The presence of short-range atomic order can have a significant

influence on the mechanical and physical properties of these

alloys. In this context, our subsequent results for alloys

containing about 1 at. % oxygen may be very significant. We have

observed that this small amount of oxygen essentially eliminates

the short-range order and replaces it with a two-phase mixture of

a and a z phases. One can expect this effect to be critically

dependent on the aluminum and oxygen contents in localized

regions of the alloy, and this may explain Shull's observation of

small a 2 phase clusters (Ref. 13).

Our direct evidence of short-range ordering in binary Ti-Al

alloys having a low interstitial content also supports Shull's

contention that discrepancies between various versions of the a +

a 2 region in the Ti-Al phase diagram can be produced by

differences in short-range ordering (Ref. 13 ). Shull pointed out

that certain types of measurement such as magnetic susceptibility

(Ref. 19) would be quite sensitive to short-range ordering while

other methods such as optical microscopy (Refs. 13 , 10 ) would be

unable to detect it at all. Thus, the position of phase
35

boundaries may differ significantly in these two types of

experiments

.

In Fig. 7 we have drawn a dashed line to separate the regions of

the aTi solid solution that are with and without short-range

ordering. This line corresponds rather well to the a solvus

boundary of Swart z endruber e t a 1 . (Ref. 19) who used magnetic

susceptibility as a primary experimental method.

C. The Ti-Al Phase Diagram up to 50 at. % A1 and above 1200°C

Ti-Al phase relationships above 1200°C are not well understood

because most previous investigators have worked with "quenched"

samples which yield ambiguous interpretations. It appears,

however, that this region of the phase diagram is ideally suited

for our h igh - t emp e r a tur e neutron diffraction methods. We have

used our new h igh - t emp e r a tur e facility to explore this region

directly for the first time. This rather preliminary study was

conducted by examining critical diffraction peaks of various

phases as the temperature was gradually raised. For example, we

looked at the strong (101) peak for the a 2 phase and the (101)

peak of aTi. (Note that these are two separate lines occurring

with a wide angular separation.) We also monitored the (110)

peak of the 0 phase. By carefully increasing the temperature in

small increments ( ~ 5 to 10°C) and waiting 10 to 15 minutes for

the sample to equilibrate between each increment, we were able to

follow phase transitions well enough to reproduce known
36

transitions in the Ti-Al system at temperatures above 1000°C.

Our high- temperature data essentially confirmed the location of

the transus and a solvus boundaries as determined in the

previous studies (Refs. 13,14,16) and are compatible with the

previously defined a 2 phase boundaries.

However, above 1200°C there were serious discrepancies. The

previous work had indicated that the ft phase region extended to

over 40 at . % A1 (Ref ,• 14) , but we failed to see any /3 pha s e lines at 3 5 or 40 at . % A1 . Instead, the pattern obtained at

about 1200 ± 10 0 C cons i s t e

d

of only one line at 2 9 = 40.6° (d =

2.22A). This line was obviously not a f3 phase line nor was it a

line of the equiatomic 7 phase. The absence of the strong (101)

superlattice line ruled out the a 2 phase.

The unidentified line had a moderate intensity and seemed to

correspond roughly to the strongest line of the disordered

hexagonal aTi solid solution. This had to be a tentative

identification, however, since none of the other lines of the aTi

pattern were be observed. The failure to see other lines might

be expected, however, since the a phase would produce a

relatively weak pattern at such high A1 contents. This is caused

by the positive and negative atomic scattering potentials of A1

and Ti atoms cancelling each other and presumably leading to zero

intensity at approximately 50 At. % A1 in a disordered Ti-Al

phase . In order to confirm our tentative identification of the
37

aTi structure, we would probably need to conduct high- temperature

x-ray diffraction studies on these samples. Such facilities were

not available to us at this time, however, so we decided to

examine a 40 at. % A1 sample that had been rapidly quenched

(cooling rate of 10 6 o C/sec) using an arc furnace equipped with a

piston and anvil device. X-ray diffraction patterns see (Fig. 8 ) obtained at room temperature from quenched foils having the

nominal composition 40 at. % A1 showed only the lines of a hexagonal (aTi) structure together with weak lines of TiAl 7

phase. No "superlattice" lines of the a 2 phase were detected.

The strongest line in the pattern occurred at d = 2.195 ± 0.01 A.

The quenched foil was subsequently annealed for a few minutes in

a high vacuum at 1100°C and its x-ray pattern was then that of

the ordered hexagonal a 2 phase, complete with "superlattice"

lines. This suggests a re-establishment of equilibrium in the

annealed sample which is compatible with the existing phase

diagram

.

It appears, therefore, that the existing phase diagrams (Refs.

13,14,20) must, once more, be modified. The diagram must

apparently include a disordered hexagonal phase stable above

1150°C whose existence had not been previously detected. One

might regard this new phase region as an extension of the aTi

terminal solid solution, but the two-phase regions are apparently

not connected in the binary Ti-Al system. We established this by

carefully heating our sample containing 30 at. % A1 in
38

temperature increments of about 3°C while monitoring its neutron

diffraction pattern for signs of the aTi phase. We found no sign

of this phase at any temperature. The strong (101) super- lattice

line of the a 2 phase remained until about 1202°C. At 1205°C the

a z phase line disappeared and was replaced by the strong (110)

line of the phase. Thus, we conclude that the two disordered

hexagonal Ti-Al phases are separate phase regions rather than a

single continuous phase in the binary system.

We have not yet acquired sufficient data to fully define the

phase boundaries in this system, but we have plotted the data

points provided by our neutron diffraction study and constructed

tentative phase boundaries (using dashed lines) that are

compatible with the data (Fig. 7). The new phase boundaries are

also compatible with previous solidus and liquidus points taken

from reference 14. Obviously, the solubility of A1 in /3Ti is

significantly less than had been previously reported, and this

reduced solubility is apparently caused by the presence of the

ne wly - d i s c ove r e d h igh - t emp e r a tur e (aTi) phase.

D. The Influence of Oxygen on the Phase Boundaries in the Ti-Al

System

Having established the binary Ti=Al phase boundaries in the

region of interest, we turned our attention to the influence of

oxygen on these boundaries. The time remaining did not permit us
O'

on

Fig.

8.

X-ray

diffraction

patterns

from

Ti

60

Al

A0

alloys

which

were

quenched

from

the

liquid

("splat"

-cooled)

compared

with

pattern

of

Ti3

Al.

Miller

indices

(hkJ?)

are

for

Ti3

Al.

Top

to

bottom

are:

(a)

Ti60

Al^0

quenched

and

then

annealed

at 1100°C,

(b)

Ti

60

Al

A0

as-quenched,

(c)

Ti3

Al.
40

to completely evaluate these influences, but we were able to

examine four alloys containing 1 at. % and 6 at. % oxygen. The

data, obtained by temperature scanning focused on the strong

(101) line of the a 2 phase, are plotted in Fig. 9 together with

tentative phase boundaries that are compatible with this data.

It appears that dissolved oxygen raises the transus and a

solvus temperatures and increases the stability of the aTi phase,

in agreement with previous work. The solubility of A1 in aTi is

apparently reduced by the presence of oxygen, in agreement with

previous results.

The most interesting effect produced by dissolved oxygen in aTi,

however, is its ability to eliminate the short-range order (SRO)

that we observed in some of our binary alloys. For example, no

SRO was seen in our alloys Ti 83 Al 16 0 and Ti 81 Al 18 Q at 1000 °C,

whereas binary Ti-Al alloys having the same nominal A1 contents

(16 and 18 at. %) both contained prominent SRO. Similarly, at

900°C we saw no SRO in our alloy Ti 83 Al 16 0 although the binary

Ti 84 Al 16 alloy showed pronounced SRO. In each case, the former

region of SRO was replaced by a region containing the two phases,

a + a 2

. Such behavior may produce significant changes in the

mechanical properties of these alloys. In general, our

preliminary results on the oxyg en - c on t a in ing alloys support the

previous studies which established oxygen as a strong stabilizer

of the a phase with respect to the ft phase. More data are

needed, however, to develop a truly quantitative picture.
Weight % Aluminum

41

0 5 10 15

Ti Atomic % Aluminum

Fig. 9. Influence of oxygen on the aTi phase region in the Ti-Al

system. Numbers adjacent to the phase boundaries signify oxygen

content (At. %). Data for alloys containing < 17 at. % Al

(unfilled triangles) are from Refs. 2 and 3. All other data

points are from this study. Unfilled circles are transus points.

Filled circles are solvus points. The location of the /? transus

boundary for alloys containing 6 at. % Al was not determined.
42

VIII. CONCLUSIONS

In order to determine the influence of interstitial elements on

the phase relationships in the Ti-Al system, one must obviously

begin with reliable information about the binary Ti-Al system

itself. We have found that the recently "assessed” phase diagram

(Refs. 14,20), based on an extensive literature dating back many

years, is still fundamentally inadequate.

The exact location of the a solvus boundary, for example, which

has been the focus of so much previous work, is still uncertain.

This is partly due to an inadequate knowledge of "coherency"

effects and partly to a lack of information about the effects of

short-range ordering in the aTi solid solution. We have found

that neutron diffraction methods can be very useful for detecting

and measuring the amount of short-range order in Ti-Al alloys.

We have also observed that the short-range ordering in these

alloys is strongly influenced by the presence of dissolved oxygen

in the amount of < 1 atomic percent. Oxygen apparently

destabilized the short-range order and replaced these singlephase regions with two-phases; disordered aTi and the long-range

ordered a 2 phase. The extent to which this "destabilization"

occurs depends sensitively on the A1 and 0 contents, and it may

have a profound influence on the strength and ductility of

commercial t it anium- aluminum alloys. More work is obviously

needed in this area.
Our high- temperature data a t about 1200 ° C revealed a new binary

Ti-Al phase at about 40 a t . % A1 which s e ems to have escaped

detection in most of the previous s tudie s b e caus e it cannot be

easily retained by quenching. The neutron diffraction pattern of

this phase contained only one peak, however, and this is

insufficient for positive identification of the crystal

structure. Rapid quenching of this alloy from the liquid state

produces a disordered hexagonal aTi structure. This

identification is compatible with the observed single neutron

diffraction peak at 1200°C if one assumes that the observed peak

is indeed the strongest peak of the aTi phase and that the weaker

peaks do not have intensities comparable to the background

scattering. The diffraction angle of the observed peak (29 = ~

40.6°) corresponds to a d-value of 2.22A, and this clearly rules

out a /3Ti or yTiAl structure but is compatible with an aTi

structure. The a 2 phase is ruled out by the absence of the

strong (101) superlattice peak.

The existence of a new Ti-Al phase at ~ 40 at. % A1 implies that

the solubility of A1 in /?Ti has been greatly overestimated in the

previous studies of this binary system (see Refs. 14,20). More

work is clearly needed, particularly high - temperature

measurements, in order to adequately define the phase

relationshps in this region.
44

We have thus far obtained only a limited amount of data from

alloys containing oxygen. It seems to indicate that oxygen has

the effect that was anticipated. The a/p transition temperatures

(/? transus) and a/a 2 boundary (a solvus) are displaced to

significantly higher temperatures by as little as 1 at. % oxygen.

More data is needed, however, to define the location of the a

solvus and a 2 phase boundaries. We had hoped to complete this

portion of our work in the allotted time period, but we have

encountered repeated and prolonged delays in obtaining permission

to purchase some critical equipment. This, combined with

unavoidable reactor "shut-downs” for maintenance and

modification, has inordinately delayed the work. Greater

flexibility in procurement procedures is needed to permit more

efficient coordination with the operating schedule of the

reactor

.

If such problems can be eliminated, one can obtain extremely

useful data from h igh - t emp e r a tur e neutron diffraction

measurements. The ability to detect short-range order and to

discriminate sensitively between the aTi and a 2 (Ti 3 Al) phases are

significant advantages of this method. Another advantage is the

ability to identify h igh - t emp e r a tur e phases that cannot be

retained by quenching, particularly if the structure involves

atomic ordering.
However, the use of neutron diffraction for simply determining

transition temperatures is inappropriate since such

determinations can be accomplished more quickly, more accurately

and less expensively with differential thermal analyses (DTA).

Perhaps the best method for determining a phase diagram, such as

Ti-Al, would be a combination of DTA, for determining transition

temperatures, and high- temperature neutron diffraction for phase

identification. The technology for conducting DTA studies on

levitated samples under high pressure of inert gases to suppress

vaporization has already existed for several years (Ref. 21) and

has been used successfully in determining the c 1 o s e ly - r e 1 a t e d NbA1 phase diagram (Ref. 22). It is time to apply these techniques

to the Ti-Al system.

We wish to conclude this study by offering our samples to other

investigators. We have prepared numerous samples of Ti-Al

ternary alloys containing small amounts of oxygen, nitrogen and

carbon in anticipation of further work. These samples have been

annealed in a very high vacuum (10' 8 Torr) for homogenization and

could be used for DTA studies as well as for neutron diffraction.

The samples should now be made available to other investigators

so that work may proceed as rapidly as possible on these vitally

important alloy systems.
46

Ac know 1 e dgment s

We would like- to express our appreciation to the Reactor group at

the National Bureau of Standards for their cooperation and for

the use of their facilities during this study. We are especially

indebted to Dr. Ted Prince for his sound advice and his generous

donations of time expended in collecting the data.

We wish to thank the Vacuum Furnace Systems Corporation for

undertaking to construct the special vacuum furnace apparatus for

heating the samples during the neutron diffraction measurements

.

The apparatus was subsequently found to be highly reliable and

capable of providing accurate data while protecting the samples

from contamination. We wish to thank Hr. Bill Lugmayer, P.E. , for his enthusiastic advice and assistance with e 1 e c t r on - b earn

welding and Mr. F. A. Schmidt of the Materials Preparation

Center, Ames Laboratory, for offering the use of their facilities

to fabricate the pure vanadium window for our vacuum furnace

apparatus. Finally, we would like to thank Mr. A. Giuseppetti

for assisting us through the many phases of sample preparation

and me t a 1 1

o

g r aph

i

c study.
IX. REFERENCES

D. Gupta and S. Weinig "The D is locat ion- Oxygen Interaction

in Alpha Titanium and its Effect on the Ductile to Brittle

Transition"

Trans. AIME 215 pp 209-216 (1959)

R . J , Van Thyne and H.D. Kessler "Influence of Oxygen,

Nitrogen and Carbon on the Phase Relationships of the Ti-Al

System"

Trans. Met. Soc. AIME 200 pp 193-199 (1954)

T.H. Schofield and A.E. Bacon "The Constitution of the

T i t an ium - r i ch Alloys of Titanium, Aluminum and Oxygen"

Journ. Inst. Met. 8 5 pp 1 9 3 - 1 9 5 ( 1 9 5 6 ) R.N. Nikitenko "Titanium and its Alloys"

Izv. Akad. Nauk SSSR p 41 (1960)

V.V. Glazova "Investigation of Phase Equilibria in the

Ternary System Titanium-Aluminum- Oxygen"

Doklady Chem: 164 894 translated from

Doklady Akad. Nauk SSSR 164 #3 p 567 (1965)

F.A. Crossley "Effects of Ternary Additions: 0, Sn, Zr, Cb , Mo and V on the a/a Ti 3 Al Boundary of Ti-Al Base Alloys"

Trans. Met. Soc. AIME 245 pp 1963-1968 (1969)

I. I. Kornilov, F.N. Tavadze, G.N. Ronami, K.M. Konstantinov,

T.A. Peradz and Yu . A . Maksimov "Influence of a- and j9-phase

stabilizing elements on the position of phase boundaries in

the Titanium- Aluminum System"

Soobshch. Akad. Nauk Gruz SSR 66 (1) pp 133-136 (1972)
48

8. V.V. Vavilova, T.A.Peradze, L.P. Fatkullina and O.S. Korobov

"Oxygen in Titanium Alloys with Aluminum"

Metal Science and Heat Treatment 1

7

#3-4, p 229 (1975)

9. J.Y. Lim, C.J. McMahon Jr., D.D. Pope and J. C. Williams

"The Effect of Oxygen on the Structure and Mechanical

Behavior of Aged Ti-8 Wt . Pet. Al"

Met. Trans. 7

A

pp 139-144 (1976)

10. F.G. Gehlin "The Crystallographic Structure of Ti 3 Al" in

"The Science, Technology and Application of Titanium" edited

by R.I. Jaffee and N.E. Promisel

Pergamon Press, New York, NY pp 349-357 (1970)

11. I . I . Korn i 1 ov , V . V . Vavilova , L.E. Fykin, R.P. Ozerov, S.P.

Solowiev and V.P. Smirnov "Neutron Diffraction Investigation

of Ordered Structures in the T itanium- Oxygen System"

Met. Trans. 1 pp 2569-2571 (1970)

12. "Phase Stability in the Ti-Al-Mo-Cb-Ta System" NBS Progress

Report submitted to David W. Taylor Ship Research and

Development Center, Annapolis, MD 21402

ONR Contract No. NO 0 1 4 - 8 C - F - 0 0 1 0 March 1 -September 30 ( 1979)

13. R.D. Shull, A.J. McAlister and R.C. Reno "Phase Equilibria

in the T i t an ium - Aluminum System" in "Titanium Science and

Technology"

Proc . Fifth International Conf. on Titanium-Munich

Vol 3 pp 1459-1466 (1984)

14. J.L. Murray "Binary Ti Alloy Phase Diagram Compilation"

ONR Final Technical Report NO 0 0 1 4 - 8 4 - F - 0 0 6 9 ( 1 98 5 )
49

15. "Phase Stability in the T i - A1 - Ho - Cb - Ta System" NBS Progress

Report submitted to ONR Materials Science Division,

Arlington, VA 22217

Oct 1 (1970) - July 1 (1980)

16. S. Sircar, K. Narasimhan and K. Mukherjee "An investigation

of the ordered D0 ig phase formation in the Ti-Al system"

Journ. Mat. Sci. 2

1

pp 4143-4146 (1986)

17. M.J. Blackburn "The Ordering Transformation in Titanium:

Aluminum Alloys Containing up to 25 at. pet. Aluminum"

Trans. Met. Soc. AIME 239 p 1200 (1967)

18. T.K.G. Namboodhiri, C.J. McMahon and H. Herman

"Decomposition of the a-Phase in Titanium- rich Ti-Al Alloys"

Met. Trans. .4 p 1 3 2 3 ( 1 9 7 3 )

19. L.J. Swar t z endrub e r , L.H. Bennett, L.K. Ives and R.D. Shull

"The Ti-Al Phase Diagram: the a-a 2 Phase Boundary"

Mat. Sci. and Eng. _51_ P1-P9 (1981)

20. J.L. Murray "Calculation of the T itanium- Aluminum Phase

Diagram"

Met. Trans. 1 9

A

p 243 ( 1 988 )

21. R. Flukiger "Phase Diagrams of Superconducting Materials"

in "Superconductor Materials Science," Chapter 8, edited by

S. Foner and B.B. Schwartz

Plenum Publishing Corp. (1981)

22. J.L. Jorda, R. Flukiger and J. Muller "A New Metallurgical

Investigation of the Niob ium - Aluminum system"

Journ. Less-Common Metals 7 5 p 227 ( 1 980 )
50

List of Tables

Table number Page

1. Manufacturers' "Typical Analyses for "High- 16

purity" Titanium (in PPM by Weight).

2. Effects of Various Melting Procedures on the 19

Analyzed 0, N and C Content of Pure "Iodide"

Titanium.

3. Analyses of Binary Ti-Al Samples After Completion 23

of All Experiments, Including Homogenization and

H igh - t emp e r a tur e Neutron Diffraction Studies.
51

List of Figures

Figure number Page

1. Schematic representation of the (aTi) 5

structure and the a 2 (Ti 3 Al) structure.

Open circles represent Ti atoms, dark

circles represent A1 atoms and shaded (grey)

circles represent a statistically random

occupancy by both Ti and A1 atoms in the

ratio called for by the phase composition.

Thick lines outline the respective unit cell

for each structure. The structure of the a 2 phase actually involves slight displacements

of the Ti atoms from these "ideal positions"

(ref . 10)

.

2. High-purity "iodide" Ti crystal bar imported 17

from the People's Republic of China; used as

the starting material of this investigation.

3 . Weight losses produced by repeated arc- 17

melting for pure iodide Ti, E 1 60 Ti and

synthetic Ti + 5% dehydrated NaCl mixture.
52

Figure number Page

4. Neutron diffraction patterns for (aTi) phase 26

(annealed Ti powder) and a 2 (Ti 3 Al) phase

(annealed Ti 3 Al). Peaks are identified by

the Miller indices of their respective unit

cells. The patterns were obtained at room

temperature using unrotated samples.

5. Schematic representation of high- temperature 28

neutron diffraction apparatus operable to

about 1500°C under a high vacuum.

6. Photograph of the completely assembled high- 29

temperature neutron diffraction apparatus

ready for attachment to the vacuum system.

Tentative revised phase diagram of the TiAl

system. This is essentially the assessed

diagram of J.L. Murray (Refs. 14,20)

modified so as to be compatible with our

observed h igh - t emp e r a tur e neutron

diffraction data. The new high- temperature

phase at ~ 40 at. % A1 is tentatively

identified as (aTi). The region of SRO is

33

also shown

.
5 3

Figure number

8. X-ray diffraction patterns from Ti 60 Al 40

alloys which were quenched from the liquid

(" splat cooled) compared with pattern of

Ti 3 Al. Miller indices (hki) are for Ti 3 Al.

Top to bottom are: (a) Ti g0 Al 40 quenched

and then annealed at 1 1 0 0

9 G , (b) Ti 60 Al^ 0 as-quenched, (c) Ti 3 Al.

9. Influence of oxygen on the qTI phase region 41

in the Ti-Al system. Numbers adjacent to

the phase boundaries signify oxygen content

(At. %). Data for alloys containing < 17

at. % A1 (unfilled triangles) are from Refs.

2 and 3. All other data points are from

this study. Unfilled circles are transus

points. Filled circles are solvus points.

The location of the /9 transus boundary for

alloys containing 6 at. % A1 was not

Page

39

determined

.
NBS-U4A (rev. 2 -so) U.S. DEPT. OF COMM. 1. PUBLICATION OR 2. Performing Organ. Report NoJ 3. Publication Date

BIBLIOGRAPHIC DATA

REPORT NO.

SHEET (See instructions) NISTIR 88-3856 OCTOBER 1988

4. TITLE AND SUBTITLE

Effect of Interstitial Elements on Phase Relationships in the

Titanium-Aluminum System

5. AUTHOR(S)

R. M. Waterstrat

60 PERFORMING ORGANIZATION (If joint or other than NBS, see instructions) 7, Contract/Grant No.

NATIONAL BUREAU OF STANDARDS

U.S. DEPARTMENT OF COMMERCE

GAITHERSBURG, MD 20899

8 . Type of Report & Period Covere

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City, State, ZIP)

American Dental Association Health Foundation

211 E. Chicago Avenue

Chicago, IL 60611

10. SUPPLEMENTARY NOTES

J Document describes a computer program; SF-I 8S, FIPS Software Summary, is attached.

11. ABSTRACT (A 200-word or less factual summary of most significant information . If document includes a si gnificant

bi bliography or literature survey, mention it here)

Neutron diffraction measurements on t itanium- aluminum alloys were

obtained at temperatures up to 1350®C in a specially constructed

high-vacuum furnace which incorporates pure vanadium as a nullmatrix material for all components which are in the path of the

neutron beam. The data Indicate that atomic short-range ordering

probably exists In alloys containing 16 to 18 atomic % aluminum

between 900 and 1000 e G. This is replaced by a two-phase (a + a 2 )

structure when the alloys contain - 1 atomic % oxygen. A new hightemperature phase was detected at ~ 40 atomic % A1 and at least

1200 e C which is tentatively identified as having a hexagonal closepacked (oTI) structure. Additions of up to 6 atomic % oxygen raise

the a solvus and f} transus temperatures of titanium- aluminum alloys.

12o KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolon s) intermetallic compounds; neutron diffraction; phase diagrams;

titanium alloys; titanium-aluminum alloys; titanium-oxygen alloys

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