2495

*“i^»oo.'vation and Substitution

Technology for Critical Materials

Volume I

Proceedings of Public Workshop sponsored by

U.S. Department of Commerce/National Bureau of Standards

and U.S. Department of the Interior/Bureau of Mines

15-17 June 1981

Vanderbilt University

Nashville, TN

April 1982

, UbG PARTMENT OF COMMERCE

G2“2^95 - bureau OF STANDARDS

V.l

1982

MAR 15 ^983

NBSIR 82-2495

CONSERVATION AND SUBSTITUTION

TECHNOLOGY FOR CRITICAL MATERIALS

VOLUME I

Proceedings of Public Workshop sponsored by

U S. Department of Commerce/National Bureau of Standards

and U.S. Department of the Interior/Bureau of Mines

15-17 June 1981

Vanderbilt University

Nashville, TN

General Chairman;

Allen G. Gray

American Society for Metals

Adjunct Professor of Metallurgy

Vanderbilt University

Workshop Coordinator;

Robert T. Nash

School of Engineering

Vanderbilt University

Nashville, TN

April 1982

U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary

NATIONAL BUREAU OF STANDARDS. Ernest Ambler, Director

TABLE OF CONTENTS

VOLUME I Page

I. Foreword F 1

II. Table of Contents TCI

Workshop Keynote Address

(1) OPPORTUNITIES IN CONSERVATION AND SUBSTITUTION TECHNOLOGY FOR

CRITICAL MATERIALS, Allen G. Gray, American Society for Metals P 1

Current Activities in Substitution

(2) TECHNICAL ACTIVITIES AT THE NATIONAL BUREAU OF STANDARDS

RELATED TO SUBSTITUTION AND CONSERVATION, John B. Wachtman, Jr.,

Center for Materials Science, National Bureau of Standards. P 2

(3) CURRENT ACTIVITIES IN SUBSTITUTION, UNITED STATES DEPARTMENT

OF INTERIOR, BUREAU OF MINES, Ralph C. .Kirby, Mineral Resources

Technology Division. P 3

(4) NASA'S ACTIVITIES IN CONSERVATION OF STRATEGIC AEROSPACE

MATERIALS, Joseph R. Stephens, NASA--Lewis Research Center. P 4

(5) CHARGE TO 'WORKSHOP IN RELATION TO MATERIALS POLICY LEGISLATION

(PL 96-479), John B. Wachtman, Jr., National Bureau of

Standards. P 5

Chromi um--Substi tuti on and Conservation

(6) CONSERVING CHROMIUM, David C. Goldberg, Westinghouse Electric

Corporation. P 6

(7) COMPUTER SYSTEMS FOR DEVELOPING CHROMIUM FREE STEELS FOR

HARDENABILITY , Dale H. Breen, International Harvester Co. P 7

(8) SUBSTITUTIONS FOR CHROMIUM IN STEELS FOR COIL SPRINGS,

S. T. Furr, Bethelehem Steel Corp. P 8'

(9) AN UPDATE ON FE-MN-AL STEELS AS A REPLACEMENT FOR NI-CR

STAINLESS STEELS, Samir K. Banerji, Foote Mineral Co. P 9

(10) SUBSTITUTE STAINLESS STEELS WITH LESS CHROMIUM,

Joseph A. Oouthett, Armco, Inc. PIO

(11) SUBSTITUTION OF MODIFIED 9CR-1M0 STEEL FOR AUSTENITIC STAINLESS

STEELS, Vinod Sikka, Oak Ridge National Laboratory. Pll

TCI
Cobal t--Substitution and Conservation

(12) COBALT, A VITAL ELEMENT IN THE AIRCRAFT ENGINE INDUSTRY,

Joseph R. Stephens, NASA--Lewis Research Center. P12

(13) REDUCTION OF COBALT IN GAS TURBINE ENGINES, Dennis J. Evans,

Pratt & Whitney Aircraft Group. P13

(14) ROLE OF COBALT IN ASTROLOY (SUPERALLOY), John K. Tien and

Robert N. Jarrett, Columbia University. P14

(15) PROGRESS IN REDUCTION OF COBALT IN MAR-M247, Michael Nathal

,

Case Western Reserve University/NASA--Lewis Research Center P15

(16) SUBSTITUTION FOR COBALT IN HIGH SPEED STEELS--ROLE OF

PROCESSING INNOVATION, Walter T. Haswell, Jr., Colt

Industries, Crucible Inc. P16

(17) SUBSTITUTION FOR COBALT IN MAGNETIC ALLOYS,

Thomas D. Schlabach, Bell Telephone Laboratories P17

(18) BORONIZING AS A SUBSTITUTE FOR COBALT BEARING ALLOYS,

Stanley R. Scales, Hughes Tool Co. P18

Titanium--Substitution and Conservation

(19) TECHNOLOGICAL OPPORTUNITIES FOR A MORE STABLE SUPPLY OF

TITANIUM, Harris M. Burte, Air Force Materials Laboratory P19

(20) ISOSTATIC PRESSING OF COMPLEX SHAPES FROM TITANIUM AND

TITANIUM ALLOYS, Stanley Abkowitz, Dynamet Technology Inc. P20

(21) BENEFITS OF TITANIUM IN AIRFRAMES AND COMPROMISES IN

SUBSTITUTION, Donald G. Goehler and Lew Wingert, Boeing

Commercial Airplane Co. P21

Tantal um--Substitution and Conservation

(22) TANTALUM AND OTHER CRITICAL ELECTRONIC MATERIALS,

Thomas D. Schlabach, Bell Telephone Laboratories. P22

(23) TANTALUM IN HIGH TEMPERATURE ALLOYS, James A. Ford and

Robert B. Herchenroeder , Cabot Corporation. P23

TC2
VOLUME II

Technological Opportunities for

Conservation of Critical Metals Utilizing

Metallurgical Coating Systems

(24) A PRESENTATION AT THE "WORKSHOP ON CONSERVATION AND SUBSTITUTION

TECHNOLOGY FOR CRITICAL MATERIAL," David S. Duvall, Pratt &

Whitney Aircraft Group. P24

(25) POTENTIAL FOR HARD FACING APPLICATION TECHNOLOGY TO CONSERVE

CRITICAL METALS, Steve J. Matthews, Cabot Corporation. P25

(26) POTENTIAL FOR CLADDINGS TO CONSERVE CRITICAL METALS,

James M. Glazebrook, Lukens Steel Co. P26

(27) CONSERVATION OF CRITICAL METALS UTILIZING SURFACE ALLOYING,

Ray J. Van Thyne, Dilex Systems P27

(28) ELECTRON BEAM IRRADIATION APPLIED TO CONSERVATION OF CRITICAL

MATERIALS, Joe E. Jenkins, Leybold-Heraeus Vacuum Systems, Inc. P28

(29) APPLICATION OF THERMAL SPRAYING PROCESSES TO CONSERVE CRITICAL

MATERIALS, Edward R. Novinski, METCO, Inc. P29

Opportunities for Conservation and Substitution of

Critical Metals Utilizing Rapid Solidification

Technology and Near Net Shape Processing

(30) OPPORTUNITIES FOR CONSERVATION AND SUBSTITUTION OF CRITICAL

METALS UTILIZING RAPID SOLIDIFICATION TECHNOLOGY AND NEAR NET

SHAPE PROCESSING, Arden L. Bement, Jr., TRW, Inc. P30

(31) POTENTIAL OF RAPID SOLIDIFICATION FOR REDUCTION OF CRITICAL

ELEMENT CONTENT OF JET ENGINE COMPONENTS, Joseph Moore and

Colin Adam, Pratt & Whitney Aircraft Group. P31

(32) POTENTIAL OF RAPID SOLIDIFICATION FOR WROUGHT P/M ALUMINUM

ALLOYS, Walter S. Cebulak, Aluminum Company of America. P32

(33) POTENTIAL FOR NEAR NET SHAPES, J. Stainley Mosier and

Peter G. Bailey, General Electric Company P33

(34) OUTLOOK FOR AS-HIP'ED NEAR NET SHAPES, John Moll, Colt

Industries, Crucible Research Center. P34

TC3
Technological Opportunities to More Fully Utilize

Metals with Domestic Sources or Near Domestic

Sources as Substitutes for Critical Metals

(35) TECHNOLOGICAL OPPORTUNITIES TO MORE FULLY UTILIZE METALS WITH

DOMESTIC SOURCES OR NEAR DOMESTIC SOURCES AS SUBSTITUTES FOR

CRITICAL METALS, Horace N. Lander, Molybdenum Co. P35

(36) MOLYBDENUM AS A SUBSTITUTE FOR CRITICAL METALS,

Ralph M. Davison, Climax Molybdenum Co. P36

• (37) NICKEL AS A SUBSTITUTE FOR METALS IN CRITICAL SUPPLY,

Robert J. Johnson, International Nickel Co. P37

(38) TUNGSTEN--A NEWLY NONCRITICAL METAL, C. C. Clark, AMAX

Tungsten. P38

(39) NEW APPLICATIONS OE VANADIUM IN STEELS TO SAVE COST, ALLOYING

AND ENERGY, P. L. Mangonon, Foote Mineral Co. P39

(40) NEW RECLAMATION PROCESS FOR CRITICAL METALS, William J. Boesch,

Special Metals, Inc. P40

Role of Composites in Substitution, Conservation

and Displacement of Critical Materials

(41) THE ROLE OF COMPOSITES IN SUBSTITUTION, CONSERVATION, AND

DISPLACEMENT OF CRITICAL MATERIALS, Alfred E. Brown,

Celanese Corp. P41

(42) ROLE OF COMPOSITES IN SUBSTITUTION, CONSERVATION, AND

DISPLACEMENT OF MATERIALS AND POTENTIAL FOR SUBSTITUTION IN

MILITARY AND COMMERCIAL AEROSPACE, Richard N. Hadcock,

Grumman Aerospace Corp. P42

(43) POTENTIAL FOR SUBSTITUTION IN ALL USES OTHER THAN AIRCRAFT?

AEROSPACE, John P. Riggs, Celanese Corp. P43

Substitution Preparedness--

Information Stockpile on Substitution Technology

(44) SUBSTITUTION PREPAREDNESS--INFORMATION STOCKPILE ON SUBSTITUTION

TECHNOLOGY, A. 0. Schaefer, Metal Properties Council. P44

(45) INFORMATION ANALYSIS CENTERS AND SUBSTITUTION TECHNOLOGY,

Harold Mindlin, Battelle Columbus Laboratories. P45

(46) INFORMATION AS AN ESSENTIAL ASPECT OF CONSERVATION,

H. David Chafe, American Society for Metals. P46

(47) FACTORS INFLUENCING A NATIONAL MATERIALS SUBSTITUTION

INFORMATION DATA BASE, Maurice A. H. Howes, ITT Research

Institute. P47

TC4
(48) INFORMATION REQUIREMENTS WHEN FACED WITH A MATERIAL SHORTAGE,

Bruce E. Boardman, Deere & Company. P48

(49) SUBSTITUTION PREPAREDNESS--THE INFORMATION STOCKPILE ON

SUBSTITUTION TECHNOLOGY, John Rumble, National Bureau of

Standards. P49

(50) REQUIREMENTS FOR AN INFORMATION STOCKPILE ON SUBSTITUTION

TECHNOLOGY, Frank M. Richmond, Universal-Cydops Specialty

Steel Div. P50

(51) SUBSTITUTION PREPAREDNESS-ROLE OF COLLEGES AND UNIVERSITIES,

Ellis D. Verink, Jr., University of Florida. P51

(52) SUMMARY AND OVERVIEW OF WORKSHOP, N. E. Promisel, Consultant

and Allen G. Gray, General Chairman. S 1

(53) SUMMARY OF FINDINGS SFl

(54) SUMM^Y OF RECOMMENDATIONS R 1

(55) PROCESS EMPLOYED TO DEVELOP THE FINDINGS AND RECOMMENDATIONS PRl

(56) WRITTEN SUBMISSION TO WORKSHOP, James I. Mueller, President,

American Ceramic Society. W 1

TC5

FOREWORD

As is noted in the Charge to Workshop (pg. P5 ) presented

by John B. Wachtman Jr. , of Department of Commerce /National

Bureau of Standards , which cosponsored the event with Department of the Interior /U. S. Bureau of Mines, this Public Workshop was held to assist in the preparation of a report to

Congress required by a law passed in October 1980, the National

Materials and Minerals Policy, Research and Development Act of

1980. As stated by Dr. Wachtman, the Act requires separate

reports from the Department of Commerce (DOC) and the Department of Defense (DOD) , both due for delivery to the Congress on

1 October 1981. The Act also assigns other responsibilities to

the Department of the Interior and to the President's Office of

Science and Technology.

The Workshop was held principally to develop information

for the report required by the Department of Commerce, but

should also be useful to the other agencies in their responsibilities. The DOC report is supposed to identify a materials

needs case related to national security, economic well-being,

and industrial productivity, to assess critical materials needs,

and to recommend programs to meet these needs.

The United States is highly vulnerable to problems in supply of critical and strategic materials and it is recognized

that there is a whole spectrum of options for responding to

such crises. While a number of supply oriented options are

under study by various groups, the focus of this Workshop was

on the technical options.

The Workshop was developed with the premise that the availability of technology for substitution, conservation and displacement is among the most viable options for reducing this country's

vulnerability to cut-off in supply of critical materials. Likewise, consideration was given to the development of the concept of

"substitution preparedness" which would provide for stockpiling

FI
2

of information on substitution and on the technology for- conservation and displacement with advanced materials.

Thus the Workshop goals were:

1. Define technical opportunities for substitution, conservation and displacement of four selected critical ( imported materials — chromium, cobalt, tantalum and

titanium.

2. Identify requirements for an "Information Stockpile."

3. Consider the potential for advanced materials, for

metallurgical coating systems, and for fuller use of

domestic or near domestic resources. Explore views

on federal actions that would be desirable to facilitate the development of these potentials.

This report includes the formal papers presented at Workshop sessions. During the Workshop information was obtained

from the participants and the attendees employing a process referred to as brainwriting. Through this procedure, written

responses to carefully developed questions were obtained for

each session during the Workshop. Immediately following the

Workshop, all the findings and candidate recommendations developed in the brainwriting process were sent to the 180 people

who attended for evaluation. The summary of findings and

recommendations from the brainwriting process is presented in

the report. Likewise, an overview/summary of the Workshop is

presented and a written submission from the President of the

American Ceramic Society is included.

Allen G. Gray

General Chairman of VJorkshop

F2
OPPORTUNITIES IN CONSERVATION AND SUBSTITUTION TECHNOLOGY

FOR CRITICAL MATERIALS

Allen 6. Gray

American Society for Materials

AMERICAN SOCIETY FOR METALS • METALS PARR, OHIO 44073

KEYNOTE ADDRESS

Opportunities in Conservation and Substitution

Technology for Critical Materials

By

Dr. Allen G. Gray

Technical Director

American Society for Metals

Metals Park, Ohio 44073

General Chairman. of Workshop and

Adjunct Professor of Metallurgy

Vanderbilt University

Presented at

Public Workshop on Conservation and

Substitution Technology for Critical Materials

Sponsored By

U.S. Department of Commerce /National Bureau of Standards

U. S. Department of the Inter ior /Bureau of Mines

Vanderbilt University

Nashville, Tennessee

15 June 1981

It's not news to those of you who are participating in this

Workshop that the United States is heavily dependent on foreign

sources for supplies of key materials that are essential to the

nation's defense and for the operation of its vital industries.

For example, this country imports almost 100 percent of the strategic aerospace metals, cobalt, chromium, columbium, tantalum and

manganese

.

Over-all the U. S. is dependent on foreign sources for 22 of

the 27 metals considered vital to this country ''s economy. The

problem affects U. S. industry over-all as much as aerospace and

potential defense needs. However, until recently there has been

all too little awareness of our enormous dependence on foreign

sources of critical nonfuel minerals.

Pl-1
2

Aside from the fact that without chromium we cannot make an

efficient durable gas turbine engine, the metal has many important

industrial applications' oil refineries, petrochemicals, conventional and nuclear power plants, tanker trucks, industrial

machinery and all stainless steels.

The essential nature of critical and strategic metals requires

that all options to provide supplies or alternatives must be

pursued. On the supply side, effective steps are being taken to

strengthen the strategic stockpile and develop domestic resources.

The development of alternatives requires that viable materials

technology options derived from subst.-' ti-'.tion , conservation, enhanced

per formance , coating systems and reclamation be available at reasonably

acceptable costs.

It should be noted that the need for high performance alloy

steels, stainless steels and superalloys requiring the strategic

aerospace metals is increasing as demands have grown for higher

durability, plus higher performance fuel efficient aircraft turbine

engines. An example is the prospect for increased requirements for

tantalum in single crystal turbine blades which will be discussed

in a Workshop session.

Likewise, it should be pointed out that a transfer of aerospace

high temperature technology will be a major asset in the development

of this country's alternative energy programs including coal gasification and liquefaction, even deep well exploitation of domestic oil

resources. It's significant that extensive tests on alloys for these

uses conducted by the Metal Properties Council have shown that a minimum chromium content of 25% is required for long term service.

Pl-2
3

This adds emphasis to the strategic importance of chromium

which is among the critical metals considered in the Workshop

Program. Chromium continues to appear more critical than the

other critical metals, and seems less secure than ever before.

Unless we develop alternative materials technologies, a cutoff

of oil supply and chromium supply would attack this country on

two fronts.

As a sidelight on this issue, I might say that in an article

in TIME magazine entitled, "Strategic Metals, Critical Choices,"

published last year, I was quoted as Technical Director, American

Society for Metals, that "A cutoff of our chromium supply could be

even more serious than a cutoff in our oil supply. We do have some

oil but almost no chromium."

The pursuit of our most important national goals dictates that

we be concerned about materials availability and that we implement

an enlightened National Materials Policy which requires the development of alternatives and contingency plans which include the technology available for substitution and conservation in the event of

a materials crisis.

Although the public is coming to understand that the U. S.

faces a serious and growing difficulty in assuring a supply of

critical materials that are essential for the nation's defense

and for the operation of its vital industries, historically the

U. S. has addressed the issue of critical materials policy only

after shortages have affected manufacturing capabilities.

The theme of this Workshop is that technology is among the

most viable options for reducing this country's vulnerability to

Pl-3
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the potential of foreign minerals cartels. Also the message is

clear that the time to start to develop these technologies for

improved materials performance and processing, substitutions, conservation, coatings systems and reclamation is now.

I believe we will all agree that the best time to develop

options to deal with our vulnerability is to decide with a cool

calculated logic what to do when we have no immediate emergency

or short supply. This is what E. F. Andrews, Vice President-,

Materials and Services, Allegheny International Inc., and a long

time supporter of a National Materials Policy, told executives

of forge companies recently at their annual meeting.

This workshop is offered as a contribution to the national

effort on the issue of materials dependency. It's encouraging that

the Reagan Administration has informed itself of the problem and

appears to be prepared to take effective steps to deal with it.

Hearings were held in the U. S. Senate earlier this year on

the National Materials and Minerals Policy Research and Development Act of 1980 (P.L. 96-479) (signed by President Carter on Oct.

21, 1980) to assess progress being made in dealing with the nation's

critical shortages of strategic materials. At these hearings one of

the questions that was raised was: What steps are required, or are

we prepared to take, if the nation were confronted with a military

or economic interruption of foreign supplies of strategically

important minerals and materials?

The objective of the Workshop is to provide input for a report

on materials needs that the Secretary of Commerce is required to

present to Congress before October 1981, as well as to point out

future directions for study and action in meeting critical materials

Pl-4
5

challenges. The report is part of the requirements specified

by the recently enacted National Materials Policy, Research and

Development Act of 1980 (P.L. 96-479), and Dr. John Wachtman of

National Bureau of Standards will comment on this specifically.

/ At this point, I would like to emphasize the importance of

developing a substitution potential for critical materials, since

I have been a strong proponent of efforts to foster the substitution initiative in testimony on the National Materials Policy Bill

in the U. S. Senate in July last year, and in the U. S. House of

Representatives shortly after the bill was introduced.

In my testimony I proposec^ that the United States should embark on an organized effort to foster and support research and

development programs that will advance the practical application

of substitution technology to reduce the impact of supply interruptions in critical metals, particularly chromium and cobalt.

Likewise, it's evident that a plan should be developed to

document known substitution technologies and "stockpile" this

information. The final session of this Workshop will be devoted

to viewpoints and discussion on what form this information stockpile

•should take to provide optimum value in "substitution preparedness."

As R & D programs are developed to plug gaps and create new

options for substituting for critical metals, this technological

information would be added to the information stockpile.

As a spokesman for "substitution preparedness," which is how

I referred to my recommendations in the Congressional testimony,

I emphasized that it is essential that appraisal and recording

of substitution procedures in specific applications be as complete

as available information allows.

Pl-5
Needs for additional information should be identified and

defined so that research and development projects can be assigned

to gather data required to expand capabilities in those areas

which appear to be sound and achievable with available means.

It should be noted that leadtime is very short in an emergency

situation. Thus a stockpile of information on substitution technology is also a stockpile of time, and a valuable yet relatively

low cost supplement to the commodity stockpile.

Substitution programs should be an important part of strategic

materials planning for individual manufacturing firms. In presentations before company management groups concerned with strategic

materials planning, I have encouraged the development of substitution programs and pointed out that the objectives of a national

initiative on substitution can be applied to internal (company) substitution R & D programs.

My recommendations for materials substitution and conservation

initiatives, included as a part of my testimony in U. S. Senate hearings last July, are summarized below:

1. Find out just how much we know about substitutes for the

most critical strategic metals, particularly chromium,

cobalt, manganese, columbium, tantalum and platinum group

metals

.

2. Appraise the reliability of the information that is available on substitutes for these critical metals and determine

what research is needed to round out available data so that

the substitute materials can be used with full confidence.

Pl-6
7

3. Document known substitution technologies and stockpile this

information in a form that can be utilized immediately by

industry and defense in an emergency. This is "Substitution

Preparedness .

4. Identify basic research programs needed to create new options

for substituting for these most critical metals and as this

information is developed, add it to the substitution information stockpile. Avoid waiting for a crisis.

5. Structure research to have substitution as its only goal.

Research priority would be determined on basis of probability

of need , and probability of success . (This enhances substitution preparedness.)

6. Develop ways to improve awareness, interaction and communications in substitution technology through workshop/conferences

as we are doing in this public workshop which includes representatives from government, industry and universities. Likewise, provide universities with information and encouragement

needed to utilize more substitution and resource availability

information in materials selection courses.

7. Emphasize research to utilize substitutes with wide distribution

and with domestic and near domestic sources such as vanadium,

,

molybdenum, tungsten and nickel. Expand research to develop

manganese-aluminum alloy systems to substitute for chromium

stainless steels on the basis of long term availability of

manganese through ocean mining.

Pl-7
8

8. Recognize that the government has an important stake in

replacements and that the nation as a whole and its security

depend upon imports of critical metals. Realize that industry is involved in solving near term problems and has increasingly turned to application type work and that if government

leaves the task of substitution research totally to industry

it probably will not get done.

9. Accept the fact that the government has justification to assume

leadership in supporting research to develop enhanced substitution and conservation potentials and in sponsoring work to

gather, store and deliver reliable information on substitutes

and conservation measures for critical materials in a form

that can be applied immediately by defense and industry in an

emergency cutoff.

Regarding the final point, although there are good reasons for

individual business firms to underwrite work on substitutes for

metals critical to their existence, I would like to comment on the

rationale for government support.

An important consideration is to realize that the tendency is

for industry not to address a critical supply situation until a material is unavailable and a crisis has developed. Then it's likely

that before anything can be done the situation would be beyond

the critical stage.

For example, as long as chromium is available and is relatively

economical, as it is today, there's no incentive to substitute for

it in heat treatable steels. But since substitution appears possible

Pl-8
9

in this instance, and we will hear more about this in the Workshop,

it makes good sense to conduct research and "stockpile" information

now on the production and heat treatment of chromium-free steels

for production of important engineering components. The data could

be drawn on immediately in a chromium emergency.

It should be noted that about 10% of chromium used annually

in this country goes into two grades of constructional alloy steels,

the 8600 series .and the 4600 series. These steels are typically

used to produce such things as gears and shafts. In this application, chromium is both highly effective and cost efficient.

However, technology can be developed through research to substitute other alloying elements for chromium. In an emergency

this would release 10% of chromium for uses where substitution is

not available, such as in jet engine alloys. This information would

be equivalent to having 10% more chromium in the stockpile available

for essential uses, and should be weighed against the cost of actually

adding 10% more chromium to the strategic stockpile.

It should be noted that new tools are available today for

studying substitution and interchangeability of alloying alements.

Computer systems provide new approaches to the design of steels

and alloys with specific properties for substitution applications

and we will hear about these techniques in the Workshop Program.

Substitution is one important approach to meeting the challenges of strategic critical metals. There is, of course, an underlying need to understand and develop other options to survive any

threatened m.etals cartel, including processing technology for jiigh

Pl-9
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performance, optimized use of material properties in design, net

shapes, reinforced plastics, utilization of domestic resources,

surface modified metals and reclamation.

Processing innovations will play a major role in advancing

both substitution and conservation technology. For instance, high

speed steels without cobalt utilizing increased quantities of

molybdenum and tungsten are being produced by hot isostatic

pressing (HIP) of powders, and we will hear about this in the

Workshop Program.

Hot isostatic pressing will make increasingly important contributions to conservation by reducing the amount of strategic raw

materials required to produce a component. Likewise, the process

enables the optimum utilization of alloying elements in design.

Rapid solidification technology (RST) holds promise for the

development of materials with higher performance as well as conservation of critical elements as v;e will hear in the Workshop.

Some of the potentials that relate to critical ma'terials, if this

new technology can be moved forward, are exciting:

1. Cobaltfree, high creep strength, nickel base superalloys

for turbine airfoils.

2. An appreciable increase in incipient melting point of nickel

base superalloys.

3. Oxidation/corrosion resistant iron-aluminum alloys and perhaps

manganese-aluminum alloys as substitutes for chromium stainless steels-.

5. Laser glazing to produce layers with different alloy

compositions

.

Pl-10
11

It also should be noted that RST promises to improve

aluminum and titanium alloys as well as superalloys as we will

hear in the Workshop.

Earlier it was noted that research efforts should be aimed

at developing materials technologies to utilize to the fullest

extent domestic and near domestic metal resources. (Incidentally,

it appears that the Soviets are giving much consideration to this.)

A workshop session is being devoted to these more available resources

including molybdenum, nickel, vanadium and tungsten. Likewise,

consideration is being given to titanium supply and improved utilization of what's available.

Also as I indicated earlier, manganese is of interest in the

long term as a resource in alloy systems utilizing aluminum to

substitute for chromium stainless steels. The long term potential

for abundant supplies of manganese from ocean mining of nodules suggests research support for alloy designers in this area. The Workshop Program will explore future directions for manganese alloy

research. In April 1978, Metal Progress published the article "An

Austenitic Stainless Steel Without Nickel and Chromium" which stirred

up a lot of interest. The author of this article is with us for a workshop presentation. In the current May issue of Metal Progress,

an article along similar lines points out that the replacement of

nickel by manganese and chromium by aluminum yields a Fe-Mn-Al alloy

that is ideal for cryogenic applications.

Enhanced activity is being generated in conservation of strategic elements by recycling and reclamation of critical metals. This

workshop will focus on a new process for reclamation of critical

metals back to quality of the virgin metals.

Pl-11
12

Application of existing and developing technology for metallurgical coating systems offers a great potential to save critical

metals. In fact, in developing the Workshop Program I arrived at

the feeling that an entire workshop could be devoted to opportunities for conservation utilizing surface modification technologies.

I will touch briefly on some of these technologies which will

be considered. Hard facing as a method of effective wear and corrosion control is receiving increasing attention to impart essential

surface properties with minimum use of chromium and cobalt. A tremendous potential exists for design of hard facing into fabrication

of original equipment in a conscientious effort to conserve critical

metals and make components last longer.

Likewise, designing with available clad steels and other clad

materials can economize on critical materials compared to use of

solid high alloy products. Clad materials properly applied can provide all the advantages of a solid critical alloy composition.

Surface modified metals offer many possible substitutes in

applications requiring solid alloys containing strategic metals like

chromium and cobalt. Several techniques are available or are in the

development stage including ion implantation, laser and electron beam

high rate fusion, spraying and surface coating with ceramics.

The concept behind surface modification is basically to put the

properties of the critical metal where they are required rather than

making the entire component from a critical metal alloy. New techniques are available for surface alloying as we will hear. High

rate surface fusion provides rapid solidification rates which give

Pl-12
13

valuable surface properties. Ion implantation with chromium or

cobalt provides the basis metal with greater corrosion and wear

resistance. Design goals can be achieved with implanted ions.

In the future we will see more composite metal components

in addition to polymer based composites which we will consider

specifically in the workshop.

We hope this workshop will stimulate the business sector to

take a more active role in dealing with the realities of critical metals.

This includes the establishment of vigorous materials management

policies and programs that cover availability, conservation, substitution, reclamation and the minimal use of critical strategic

materials

.

Critical material requirements should be considered in new

product and process developments and in long term planning. Specific

technical responses should include research in cooperation with

customers and suppliers which focuses on systems based on available

metals and alloys with only small additions of critical materials.

Also, work should be pushed to enhance the mechanical properties of

less critical alloys in an effort to meet performance requirements for

alloys containing strategic elements as well as to explore the use of

advanced materials and coating systems which will displace or reduce

the over-all need for critical metals.

Finally, it is most important that this type of information on

critical materials issues be discussed openly at all levels of the

company and especially at corporate levels where the discipline of

assignments can be established.

Pl-13
14

As I said earlier this is a "technology" workshop with the

message that technology can provide answers (some say the best

answers) for reducing this country's vulnerability to the potential

of cutoff of supplies of critical metals.

An important reality, however, is that technology cannot provide quick fixes in most cases. It takes time and considerable investment to develop, quantify and establish designer confidence in a new material or even a modified alloy. These observations were among

those made by Dr. Arden Bement , a participant in the workshop, in

a talk "Strategic Minerals - A Resource Crisis" before the New York

Chamber of Commerce and Industry and Council on Economics and

National Security. i However, Dr. Bement went on to conclude, "An overriding reality

is that technological solutions will come much faster than the

political and social solutions necessary for the peaceful international sharing of the finite resources of our planet."

One final note which helps to emphasize the objectives we have

set for the workshop comes from an address I heard in New York about

two weeks ago. The occasion was the annual meeting of the American

Iron and Steel Institute and the speaker was Eugene A. March, Group

Vice President, Colt Crucible Industries Inc., on the subject,

"Needed: A National Strategic Metals Policy."

Mr. March stated: "The severity of the critical materials

problem is such that we must tackle it on both fronts: Decrease

demand to the extent we can through substitution and conservation,

and increase the supply to the extent we can through exploration,

research and the astute practice of international politics."

Pl-14
15

I believe we will all agree that this workshop will provide a better understanding and point out future directions for study and

action on one of the fronts deemed so important by Mr. March.

Let me say that in behalf of the sponsors, the U. S. Department

of Commerce/National Bureau of Standards and the U. S. Department

of the Inter ior /Bureau of Mines, we sincerely appreciate your participation in this important program to provide input to the materials

policy direction as well as point out future directions for study

and actions.

As I indicated earlier. Dr. Wachtman will talk about this and

I am sure we will all look forward to Mr. Milbergs' comments on "Developing a Policy Framework on Materials Issues" at the workshop

dinner

Pl-15
TLCHNICAL ACTIVITIES AT THE NATIONAL BUREAU OF STANDARDS

RELATED TO SUBSTITUTION AND CONSERVATION

J. B. Wachtman

Center for Materials Science

P2
Technical Activities at the National Bureau of Standards

Related to Substitution and Conservation

Workshop on Conservation and Substitution Technology

for Critical Materials

Vanderbilt University

June 17, 1981

John B. Wachtman, Jr.

Director, Center for Materials Science

National Bureau of Standards

The Center for Materials Science carries out a comprehensive materials

research program of about $23 million per year including about $8 million per

year received from other Federal agencies for technical support of their programs.

Slide 1 (P2-5). This work relates to processing, properties, and durability of engineering materials. In addition, it addresses nondestructive evaluation throughout the processing and use of engineering materials. I cannot

possibly give a complete overview in fifteen minutes, but I would like to give

you a feeling for the character of our program and some examples of work relating

to substitution for critical materials or conservation of critical materials.

Slide 2 (P2-6). Our overall program has an impact on critical materials

in two principal ways. First, it assists in the development of new, advanced

materials which can displace critical materials in some applications. Second,

it assists in the conservation of conventional materials through helping to

reduce processing losses, improve design, and extend lifetime.

Slide 3 (P2-7). The special character of our work is its orientation toward

measurement technology, accurate data, and physical models needed for effective

utilization of materials. Let me give an example of our work in each of these

three areas.

Slide 4 (P2-8). In the area of measurement technology we are developing

improved measurement methods to assist in safe, but not excessively conservative

design with structure, containers, etc., involving alloys containing critical

materials. Shown here is a procedure for measuring the value of the J-integral,

a measure of toughness useful in design based on elastic-plastic fracture

mechanics. Current work here is focused on pipeline steels, pressure vessel

steels, ship steels, and steels for cryogenic applications.

Slide 5 (P2-8). In the area of data, we measure and provide many types.

Also, we operate a number of data centers which evaluate data from the literature and in some cases provide calculated estimates. One which is important to

development of substitute alloys is the Alloy Phase Diagram Center operated

in conjunction with the American Society for Metals.

Slide 6 (P2-10). We are currently editing ASM's quarterly, Bui 1 eti

n

of

A1 1 oy Phase Pi agrams , and developing computer programs to aid in evaluation,

plotting, and display of diagrams. Shown here is a display of one section of

'the iron-chromium-nickel diagram.

P2-1
Slide 7 (P2-11). Turning to the area of basic understanding, expressed

in the form of models to support development of improved materials or improved

use of materials, consider this example of a fracture map for silicon carbide

which is a candidate material for replacing superalloys at high temperatures in some applications. This map, built up by piecing together individual theories

and experiments shows that five different types of mechanical behavior take

place in different regions of the space of temperature and stress intensity

factor.

Slide 8 (P2-12). Additional examples of materials substitution or displacement can depend on advances in processing or advances in reliability.

Let us take one example from metals, polymeric composites, and ceramics.

Slide 9 (P2-13). We are going to hear a good deal later in this meeting

about the potential of rapidly solidified alloys to replace critical materials

in some applications. The properties of rapidly solidified alloys depend on the special structures and microstructures which can be developed. We are attempting to improve conventional theory and check it by experiment to

trace out regions of concentration and velocity favoring particular microstructures. Hopefully, this^will have predictive power and aid in production

of improved materials.

Slide 10 (P2-14). Organic-matrix composites are also a very promising

material for greater use and we are going to hear more about them later in the

Workshop. Problems with delamination, especially during impact loading do exist

One approach to reduce this problem through the use of rubber-modified epoxy

resins is under study in our laboratory and at the Naval Research Laboratory.

In these materials the rubber is present as tiny particle inclusions in the

epoxy matrix. Slide 11 (P2-15). A localized stress-rel i ef zone is produced

which blunts cracks and flaws in the material. Slide 12 (P2-16). Fracture

energy values have been measured on the original resin (205) and the rubbermodified resin (F-185). These show large increases in fracture energy both

for the neat resin and in composites.

Slide 13 (P2-17). Concerning the field of ceramics as substitutes, I have

already mentioned silicon carbide and a fracture map representation of its

different regions of mechanical behavior. Different types of safe-life

design methodology can be used for these regions. Here is a design diagram

for regions of slow crack propagation. Using the method summarized by this

diagram, a minimum safe life can be assured for each combination of proof

stress and operating stress.

Slide 14 (2-18). Many uses of critical materials as alloying elements are

to enhance the corrosion-, fracture-, or wear-resistance of engineering

materials. We have work on conservation by increasing durability of all

these types while using less critical materials. J_et us take one example

concerning corrosion and two examples concerning wear.

Slide 15 (P2-19). A new, low-noise instrument developed at NBS allows one

to measure the frequency spectrum of microvolt level fluctuations in voltage

which we have found to correlate well with localized corrosion; i.e.,

pitting. Here are results on a new alloy Fe^^Ni^rCr,

.

8 ^ 26 ^ produced by

rapid solidification. When in crystalline form TocaTizea corrosion occurs

P2-2
unless an applied voltage is raised to a critical level. The same composition, when quenched to a glassy form, does not undergo localized

corrosion even under conditions causing localized corrosion in crystalline

material

.

Slide 16 (P2-20). Here are results showing how nickel-phosphorus coatings

on 0.2 tool steel can be varied to give much lower wear rates than either

uncoated tool steel or a conventional coating. The lowest rate of dry

sliding wear is produced by an electroless coating technique but this is a relatively expensive procedure. Conventional electroplating produces a relatively soft coating. By careful adjustment of the electroplating

conditions a very hard, wear-resistant coating can be produced. It is

particularly interesting to notice that pulsed electroplating gives a harder coating than conventional direct current electroplating, both in

the as-plated and heat-treated conditions.

Slide 17 (P2-21). Our final example concerns another aspect of wear resistance. There are two-phase alloys with a hard carbide phase. For certain

conditions of dry abrasive wear, the alloys with coarse carbides are much more wear resistant at the same hardness value of the alloy.

P2-3
NBS MATERIALS SCIENCE

• Processing Science

• Non-Destructive Evaluation

• Structure Properties Science

• Durability Science

P2-4
NB$ PROGRAM IMPACT IN CRITICAL MATERIALS

Measurement and Science Base for:

• Development of Advanced Materials

• Conservation of Conventional Materials
NB$ PROGRAM OUTPUT

Measurement Technology

Essential Data

Basic Understanding
LOAD

P2-7

LOAD
NB$ MATERIALS DATA CENTERS

• Diffusion Data for Metallic Alloys

• Critically Evaluated^ Computerized Alloy Phase

Diagrams

• Phase Diagrams for Ceramics Processors and Users

P2-8
P2-9
SILICON

CARBIDE

(NC

203)

IN

AIR

h- cDiO^rocj ”“0

P2-10

500

1000

1500

TEMPERATURE

(°C)
MATERIALS SUBSTITUTION & DI$PLACEf€NT

• Metals

• Organic Composites

• Ceramics

P2-11
EFFECTS

OF

RAPID

SOLIDIFICATION

ON

ALLOY

HOMOGENEITY

o/« N0liVaiN30N03 AOTTV 90T

P2-12

LOG

INTERFACE

VELOCITY

cm/s
P2-13
P2-14
>

o

cc

LU

LU

LU

CC

D

K

O

<

CC

'id-

(/) —J (U D 0.

Q O

O

00 CNl cJ cvi

CO o

1-^ ^

CO CM

o o

<3^ 0> <3^ in cn

CC

LU

H

<

Z

z if)

if) LU

LU

CC

CC in

00

in

o 1 CM LJ.

^ i o d

O (/>

(/) <

if)

if) < O

O in

00 m •,- o I CM ll

P2-15
• ‘3WH1 WnilVd MOWINIM

P2-16

Figure

18.

Design

diagram

based

on

proof

testing:

silicon

carbide,

1400*C

(Wlederhorn

et

al,

ref.

95).
MATERIALS CONSERVATION

Corrosion

Fracture

Wear

Metals

Ceramics

Polymers
CURRENT

AMPLITUDE,

A/cm'^

VHz

FREQUENCY, Hz

P2-18
WEAR

RATE

(10

mm

/m)

3

Fig. K. Observed relation between wear rate (ran /m) and Knoop liardness

(kg/mm^) for the coatings studied. For the lii-P coatings, the

wear rate shows a dependence on H

P2-19
P2-20
CURRENT ACTIVITIES IN SUBSTITUTION

UNITED STATES DEPARTMENT OF INTERIOR. BUREAU OF MINES

Ralph C. Kirby

Minerals Resources Technology Division

P3

CURRENT ACTIVITIES IN SUBSTITUTION

UNITED STATES DEPARTMENT OF THE INTERIOR

BUREAU OF MINES

Ralph C, KlL'by

Director

Mineral Resources Technology Division

Presented at

Workshop on Conservation and Substitution Technology

for Critical Materials

June 15-17, 1981

at

Vanderbilt University

Nashville, Tennessee

P3-1
1

For seven decades the Bureau of Mines of the Department of the Interior has

been the Federal Government's primary agency for research in minerals and

mineral raw materials. Its mission, Figure 1, is to help ensure the continued

viability of the domestic minerals and materials economy, and to help ensure the

maintenance of an adequate minerals base for the United States.

The fundamental program for minerals and materials R&D resides in the

Mineral Resources Technology program. It is aimed at anticipating and providing

solutions to the technological aspects of national minerals problems; helping to

ensure the availability of the critical and strategic mineral resources upon

which our industrial economy is built. The program. Figure 2, should help:

* provide technology for an adequate supply of minerals for the

United States,

* for conserving and making the best use of what we have,

* and if we need it and don't have it, doing something to get it.

The Mineral Resources Technology program addresses the functions that span

the entire Bureau of Mines portion of the mineral cycle except land rehabilitation, as shown in Figure 3. One usually associates the Bureau's R&D efforts

with mineral supply problems or the raw materials end of the mineral cycle.

However, the Bureau has made many contributions to materials development, and to

materials recycling.

Although this conference is on conservation and substitution for critical

materials , the Bureau of Mines uses a broader definition of substitution that

includes the substitution of one resource for another. In certain cases some

resources can serve as substitutes or alternatives for others. The substitution

of nonmagnetic taconites for the "direct-shipping" iron ores is one example.

P3-2
2

In the late 1950’ s, the Bureau’s analysis of the domestic iron ore situation, Figure 4, suggested that technology to process non-raagnetic taconites

would be more than desirable in 15 or so years. Bear in mind, of course, the

import picture suggested then does not obtain now. In fact, in 1980, we imported only 22% of our iron ore needs. But to help solve the predicted problem, we

started pioneering work to concentrate fine-grained hematites in 1957.

As the high-quality "direct-shipping" iron ore was depleted from the great

iron ranges of Michigan and Minnesota, magnetic separation technology to enrich

the lower-grade magnetic ores was developed.

However, vast tonnages of nonmagnetic ores lay untouched and worthless. In

the laboratory, and then in the pilot plant shown in Figure 5, the Bureau of

Mines developed the technology to recover the nonmagnetic ore. The problem was

that, when the ore was ground finely enough to liberate the iron, too much iron

was rejected with the waste. We developed technology to selectively flocculate

or clump the fine particles of iron together so that they would not be carried

away with the waste.

Based on the patents that resulted from Bureau research and on cooperative

development, the Tilden plant started production in 1974 at a capital cost of

$200 million, employing 600 people. Through the array of 150 flotation cells

shown in Figure 6, each with 500 cubic foot capacity, passed material that made

4 million tons of iron ore pellets per year. The capacity of the plant has

since been doubled to produce 8 million tons of pellets from 20 million tons of

ore—about 10% of U.S. iron ore output.

Between 1957 and 1970, the Bureau spent about $2 million for research on

the jasperite ores of upper Michigan. Staff analysis reports the value to the

economy of the period 1972-1980 was approximately $1.5 billion.

P3-3
3

The elements of the Mineral Resources Technology program (Figure 7) cover

both materials supply and materials utilization. This program is concerned with

mining more efficiently, extracting more in processing plants, and converting

domestic mineral resources to reserves. It addresses two national problems of

major concern: the technology needed for an adequate supply of minerals and

materials, and the need to improve productivity.

This year, this program is funded at roughly $26 million. Figure 8, with

the largest part going to the sub-program on Mineral Science and Technology.

The budget proposed by the previous Administration allowed a 6.8% increase for

« inflation. The program itself had an increase of $4,365 million, of which

$900,000 of this was mandated for energy conservation steps at two research

centers. The remaining $3,465 million is distributed as follows (Figure 9):

* $1 million is requested to support basic research in minerals crushing

and grinding, fine particle processing fundamentals, and thermodynamic

and kinetic aspects of cobalt reactions.

* $2 million is requested for new materials for energy development — Research to develop materials that provide adequate, not best,

performance with a minimum of critical metals in composition — Research that will lead to materials adequate for the mining and

processing of new domestic energy resources and that can reduce

critical materials needs in the equipment and facilities used in

the nonfuel minerals industry, particularly for mining and

processing equipment.

* $0,465 million is requested for solvent extraction and ion exchange of

cobalt and to reassess low-grade onshore manganese resources.

P3-4
In recent weeks the media (Figure 10) has picked up the warning from business and government leaders regarding our heavy dependence on imported strategic

minerals. We propose to Increase our effort on this important national problem.

But, it should be pointed out that of the 19 materials listed as critical

by the Council on International Economic Policy in 1974, Figure 11, the Mineral

Resources Technology program is presently conducting research on all but two

mercury and natural rubber. And we've worked on improved mercury processing in

the recent past.

Specific proposals for research on critical minerals have resulted from

reprogramming of funds by Secretary Watt from the Minerals Environmental Technology Program. This reprogramming , (Figure 12) reflects the new Administration's

emphasis on Strategic and Critical Minerals and reflects Instructions in

PL 96-479 (National Materials and Minerals Policy, Research and Development Act

of 1980) to "...increase the level of mining and metallurgical research by Bureau

of Mines in critical and strategic minerals."

In essence. President Reagan's revision of the 1982 budget proposes a 50%

increase for the Mineral Resources Technology Program. Critical and strategic

minerals research is distributed among all three subprograms. Figure 13.

Advancing Mineral Science and Technology is the largest subprogram, and

includes basic research and all extraction technologies. It covers a broad

portion of the technological innovation spectrum—from concept to pilot plant.

The following are examples of some of the research on critical materials

conducted under this subprogram.

The United States depends on Imports for most of its cobalt and nickel

supply. However, about 10% of our nickel is smelted as ferronickel from one

P3-5
deposit in Oregon. There are numerous other nickel-cobalt laterite deposits in

southern Oregon and northern California of somewhat lower grade, Figure 14.

Using a novel Bureau-developed process (Figure 15) to treat these laterites,

we not only recover nickel in an environmentally acceptable way, but we also

recover cobalt as well. The laterite deposits also contain 2% chromium; we are

developing the technology for recovering it from the residues after the nickel •

and cobalt have been leached. The new technology is being perfected in a contract pilot plant in Tucson, Arizona, operated for us by UOP, Inc. (Figure 16).

The largest known cobalt reserve in the United States is in the Blackbird

Mountain area of Idaho. It could supply as much as 20% of our cobalt needs for

15 years. Although the deposit has been previously worked, production ceased

because of severe metallurgical problems. To help solve these problems, we are

working on leaching-electrowinning techniques to recover the cobalt. Copper,

present in the ore is also recovered. The ore is leached under pressure to put

the cobalt into solution, without releasing arsenic to the environment. Once

in solution the metallic copper and cobalt can be extracted by electrowinning

(Figure 17).

The sub-program titled Conserving Domestic Mineral Resources is designed

to help conserve and make the best use of those mineral resources that we have.

It is aimed at getting all mineral values from domestic resources, and at

making use of bypassed ore bodies that were initially uneconomical to recover.

Recycling is also a major part of this subprogram. The following three projects

illustrate the scope of our recycling program with respect to critical materials

such as chromium. Our research efforts cover the chemical, metallurgical, and

refractories aspects of chromium recycling.

P3-6
Chromic acid etching solutions are used for brass finishing, etching of

printed circuits, anodizing, and preparation of plastics for plating. *The

solutions are discarded when the etching rate becomes too slow due to chemical

reduction of chromic ions and/or contamination of the solution. The Bureau

system, (Figure 18), regenerates the solution by electrochemically oxidizing

the chromium ions and removing contaminants. Three industrial plants have

tested the unit. One has already constructed their own unit. A fourth company

is in the process of constructing their own unit.

Research has developed technology to recycle stainless steelmaking particulate wastes such as electric furnace dust, mill scale, grinding swarf, and AOD

dust. The waste dusts are pelletized with a small quantity of portland cement

before introduction into the furnace. Ten tons of pelletized waste are being

introduced into the stainless and specialty steelmaking furnace. Figure 19, to

determine melting conditions.

Furnace operations such as shown in Figure 19, require refractories, most

of them containing some chromite. The Bureau is investigating methods to

recover chromium-bearing refractory grains from used furnace refractories.

Figure 20, and reusing them.

A major element of the last of our three subprograms. Developing Domestic

Mineral Resources, focusses on both alternative mineral resources and substitute materials.

Searles Lake, California, is a large desert lake in San Bernardino County

about 150 miles northeast of Los Angeles. It contains very dilute dissolved

tungsten, about 55 ppm in highest concentration. The large volume of lake brine

makes Searles Lake a significant resource in relation to current U.S. tungsten

reserves. Searles Lake brines are too dilute to process alone for tungsten, but

P3-7
7

it could be recovered as a byproduct during commercial processing that produces

potassium salts, soda ash, salt cake, and boron chemicals. We are working

cooperatively on a cost-sharing basis with the Kerr-McGee Chemical Corporation,

which operates the plant.

The key step in the recovery technology is an ion-exchange resin that the

Bureau developed specifically for these brines. When the equipment (Figure 21)

is scaled up to process the full 1,800 gal/min feed stream of the West end

plant, tungsten recovery will be equivalent to about 2 percent of current U.S.

consumption. Additional tungsten from the larger Argus plant could increase

that figure to 7%.

The shaded area of Figure 22 shows the relative portion of Aluminum raw

material , that is bauxite and alumina, that we import. Looking at this another

way, the United States builds some of the most technologically advanced military

hardware in the world, such as the F-14 fighter plane shown in Figure 23-A.' The

airframe of the F-14 is made primarily of aluminum. Where other specific

material properties are required, steel, titanium, and boron composites are

used. Only about 7%—just barely the equivalent of the vertical stabilizers—of

the airframe is built from home grown alumina and bauxite (compare the shaded

areas of Figures 23-B and 23-C).

To reduce our dependence on imported bauxite as a source of alumina, we are

obtaining process technology information from our small pilot or miniplant

(Figure 24) in Boulder City, Nevada. Here, a variety of processes for

recovering alumina from domestic nonbauxitic resources, such as clay and coal

wastes, are being tested. Five alumina-producing companies (Figure 25) are

joined with the Bureau in this work on a cooperative, cost-sharing basis.

P3-8
8

Improved fabrication methods can reduce the quantity of raw materials

needed. About 14,000 lbs. of critical materials are needed to build the engine

shown in Figure 26. The fly-away or final weight of the engine is only about

3,000 lbs. The difference is what is lost in the manufacturing processes and

machining of parts. For example, a forged titanium hub for a helicopter rotor

blade weighs about 1,000 lbs. before machining and 150 lbs. after. The

buy-to-fly ratio is about 6.6 to 1. Sikorsky Aircraft Division of United

Technologies Corporation believes that through manufacturing research on casting

technology the buy-to-fly ratio can be reduced to better than 2:1.

Powder metallurgy technology is another way to reduce the buy-to-fly ratio.

Powder can be pressed into most any shape desired. Figure 27. It can then be

sintered to make it a single coherent part. There is minimal waste this way from

machining, because after sintering the part has nearly its final shape.

The main thrust of our powder metallurgy effort is directed toward obtaining

high purity reactive metal powders. Powder metallurgy components of titanium and

titanium alloys could be possible substitutes for nickel-and cobalt-based

superalloys in low temperature applications. These superalloys are critical to

national defense, and yet nickel and cobalt ores are largely imported, whereas

the titanium ore, ilraenite, is domestically abundant. Titanium could substitute

for some superalloys if titanium powders could be made free of chloride and

tungsten impurities, which significantly affect weldability.

Bureau research has shown zinc is a possible substitute for lead-tin and

for silver solders in certain applications. More important, however, is the

ability of zinc solders to fill a temperature gap. Figure 28, between the

melting point of silver solder and the melting point of lead-tin solders. The

Bureau's zinc soldering system works well with both aluminum and copper. When

P3-9
9

aluminum heat exchangers such as shown in Figure 29 fail in service they are

generally discarded as scrap. Old field repair techniques opened more seams

than they repaired on aluminum units. The new soldering system makes it possible

to repair these units, extending service life. In addition, the higher melting

points of the zinc solders permit radiators. Figure 29, to be operated at

higher temperatures. Therefore, they can be made smaller, using less metal, and

engine efficiency is improved.

Corrosion-resistant steels can be made by substituting elements such as

aluminum and molybdenum for the critical material chromium. Figure 30 shows

the corrosion resistance of one of the new chromium-free alloys as compared to

that of ordinary low carbon steel. Both were oxidized for the same time at

750“C. The new Fe-8Al-6Mo-0. 8Zr-0. 1C alloy steel is a potential substitute for

stainless steels containing critical materials in some applications.

Outreach to industrial, academic and other governmental sectors involved

with conservation and substitution of critical materials to obtain input for

programming decisions is an important and continuing part of management of the

Mineral Resources Technology program. Many of our projects, such as those listed

in Figure 31 are done in cooperation with members of the industrial sector.

Communication ‘With members of the minerals industry is important to our program.

The input is an important segment of our program planning.

P3-10
10

BUREAU OF MINES

MISSION

ENSURE THE CONTINUED VIABILITY OF THE

DOMESTIC MINERALS AND MATERIALS

ECONOMY AND THE MAINTENANCE OF AN

ADEQUATE MINERALS BASE, SO THAT THE

NATION’S ECONOMIC, SOCIAL, STRATEGIC

AND ENVIRONMENTAL NEEDS CAN BE

BETTER SERVED

Figure 1. Mission of the Bureau of Mines of the U.S. Department of the Interior

MINERAL RESOURCES

TECHNOLOGY PROGRAM

• PROVIDES TECHNOLOGY WHICH CAN MAINTAIN

AN ADEQUATE SUPPLY OF MINERALS FOR THE

UNITED STATES

• PROVIDES TECHNOLOGY WHICH CAN

ENCOURAGE CONSERVATION AND BEST USE OF

DOMESTIC MINERAL SOURCES

• PROVIDES TECHNOLOGY WHICH CAN EXTEND THE

LIFE SPAN OF DOMESTIC MINERAL RESERVES AND

ENCOURAGE SUBSTITUTION OF ABUNDANT

MINERALS FOR THOSE WHICH MUST BE IMPORTED

Figure 2. Mineral Resources Technology Program

P3-11
BUREAU OF MINES ROLE

IN THE MINERAL CYCLE

11

BUREAU OF MINES ROLE

MINERAL RESOURCES TECHNOLOGY PROGRAM

Figure 3. Mineral Cycle Showing Bureau of Mines and Mineral Resources

Technology Program Roles

DOMESTIC IRON SUPPLY/DEMANO TO YEAR 2000

Figure 4. Bureau Analysis of Domestic Iron Ore Situation (1955-2000)

P3-12
12

Figure 5. Bureau of Mines Pilot Plant to Recover Nonmagnetic Ore

Figure 6. Bureau of Mines Research and Cooperative Development

at Work at Tilden Plant

P3-13
MINERAL RESOURCES

TECHNOLOGY

13

• Advancing Mineral Science and Technology — Mineral Science Base — Efficient Extraction Technology • Conserving Domestic Mineral Resources — Maximum Resource Recovery — Scrap and Waste Utilization • Developing Domestic Mineral Resources — Use of Plentiful Resources — Substitutes for Critical and

Strategic Minerals

Figure 7. Elements of Mineral Resources Technology Program

FY 82 BUDGET ESTIMATES

($000)

CV

FY 82

r Y O I Current Change Enhanced

Advancing Mineral Sci.

and Technology $14,143 $15,021 1

[ + $1,000

I 1+ 900) $16,921

Conserving Domestic

Mineral Resources 5,885 6,283 — 6,283

Developing Domestic

Mineral Resources 6,192 6,656

I

[+ 2,000)

[+ 4651

9,121

$26,190 $27,960 $4,365 $32,325

Figure 8. Mineral Resources Technology 1982 Program Funding

P3-14
FY 1982 PROPOSED INCREASES 14

• Basic Research, $1.0 Million — Crushing and Grinding — Fine Particle Processing — Cobalt Chemistry

• Materials, $2.0 Million — Wear- and Corrosion-Resistant Materials — Erosion/Abrasion-Resistant Materials — High-Temperature, High-Pressure Equipment

• Extractive Technology, $0,465 Million — Solvent Extraction of Cobalt — Reassess Technology for Manganese Recovery

from Low-Grade Onshore Resources

Figure 9. Distribution of Proposed Increases in FY 1982 Budget

Hov/

Figure 10. Media Focus on U.S. Dependence on Imported Strategic Minerals

P3-15
CRITICAL MATERIALS

U.S. Council on International Economic Policy, 1974

15

Aluminum • Zinc • Columbium

Chromium • Tin • Vanadium

Platinum • Titanium • Fluorspar

Iron Ore • Cobalt • Copper

Nickel • Mercury • Phosphate

Natural Rubber • Tungsten

Manganese • Lead

Figure 11. Critical Materials Determined by Council on International

Economic Policy (19174)

REPROGRAMMING

($000)

PL 96-479: . .increase the level of mining and metallurgical research by

Bureau of Mines in critical and strategic minerals. . ."

• Increase, $8,600 — Chromium . .

... $2,550 — Physical Chemistry . . $1,000 — Cobalt ... 1,450 — Mining 500 — Manganese.. ... 750 — Particulate Mineralogy . .

.. 500 — Titanium . . .

. . . 700 — Process Water 400 — Zinc . . . 250 — Scrap 300 — Tin ... 200

Figure 12. Reagan Administration Budget Reprogramming Reflects

Public Law 96-479, National Materials and Minerals

Policy, Research and Development Act of 1980

P3-16
16

FY 82 BUDGET REVISIONS

($000)

REQUEST

PENDING

PROPOSED

AMENDMENT

REVISED

REQUEST

Advancing Mineral Sci.

and Technology $16,921 +$4,100 $21,021

Conserving Domestic

Mineral Resources 6,283 + 1,700 7,983

Developing Domestic

Mineral Resources 9,121 + 2,800 11,921

$32,325 $8,600 $40,925

Figure 13. Distribution of FY 1982 Budget Revisions Among Mineral

Resources Technology Subprograms

P3-17
LATERITE ORE

COBALT

NICKEL

COPPER

Figure 15. Bureau-Developed Process for Treating Laterites Recovers

Nickel as Well as Cobalt

Figure 16. Pilot Plant for Recovering Nickel and Cobalt

is Operated by UOP, Inc.

P3-18
18

Figure 17. Extraction of Metallic Copper and Cobalt by Electrowinning

Figure 18. Bureau System for Regenerating Acid Etching Solutions by Electrochemically Oxidizing Chromium Ions and Removing Contaminants

P3-19
19

Figure 19. Steelraaklng Furance Used to Determine Melting Conditions

of Pelletized Stainless Steelmaking Waste Dusts

Figure 20. Bureau Effort to Recover Chromium-Bearing Refractory Grains

From Used Furnace Refractories for Reuse

P3-20
20

Figure 21. Equipment Used to Recover Tungsten j:. •'m Searles Lake Brines

Key to this effort is the Bureau-develo^ 'd ion-exchange resin.

U.S. SOURCES AND USES OF

ALUMINUM

DOMESTIC MINES BNET IMPORTS

Figure 22. U.S. Sources and Uses of Aluminum

P3-21
21

Figure 23-B.

F-14 Airframe With Critical

Materials Composition Shown

Figure 23-C.

F-14 Showing Home Grown

(Domestic) Alumina and

Bauxite Composition

P3-22
22

Figure 24. Alumina Miniplant in Boulder City, Nevada

PARTICIPANTS IN ALUMINA

MINIPLANT PROJECT

ALUMINUM COMPANY OF CANADA, LTD.

ANACONDA ALUMINUM

BILLITON INTERNATIONAL METALS

KAISER ALUMINUM AND CHEMICAL CORP.

REYNOLDS METALS COMPANY

Figure 25. Industry Cooperators in Bureau Alumina Effort

P3-23
23

Figure 26.

Figure 27.

TITANIUM 5,366 LB

CHROMIUM 1,656 LB

ALUMINUM 720 LB

TANTALUM 3 LB

NICKEL 5,204 LB

COBALT 910 LB

COLUMBIUM 171 LB

Weight of Critical Materials Needed to Build 3,000-lb. Engine

Powder Metallurgy Technology Yields Powder That Can be Pressed

Into Most Shapes

P3-24
24

SOLDER/BRAZE JOINING GAP

^o'to Oewofch C*ntfr

Figure 28. Solder/Braze Joining Gap

Figure 29. New Soldering System Permits Repairs to Aluminum Units,

Extends Service Life

P3-25
Figure 30. Corrosion Resistance of New Cnromium-Free Alloy Compared to

That of Ordinary Low Carbon Steel

INDUSTRIAL COOPERATION

Raise Hoist Homestake Mining Co.

Lead Project 4 Lead Producers

Iron Ore MIFGA, AISI

Phosphate TFI, Florida Inst, of

Phosphate Res.

Retractable Bit Longyear Co.

Borehole Mining Mobil, Agrico

Chromic Acid Bell Helicopter,

Keeler Brass

Recycling NARI, AFS

Tungsten Kerr-McGee

Sulfur Concrete ASARCO, The Sulphur Inst. Alumina Miniplant 5 Alumina Producers

Corrosion Natl. Assn. Corrosion Engrs.

Critical Materials ASM

Figure 31. Industrial Cooperators to Bureau of Mines Projects

P3-26
NASA'S ACTIVITIES IN CONSERVATION OF STRATEGIC AEROSPACE ,MATERIALS

Joseph R. Stephens

NASA-Lewis Research Center

P4

NASA’S ACTIVITIES IN THE CONSERVATION OF STRATEGIC AEROSPACE MATERIALS

Joseph R. Stephens

Lewis Research Center

SUMMARY

NASA has several activities underway directed at conserving strategic

materials used in the aerospace industry. Research efforts involving universities and industry as well as in-house activities at the NASA Lewis

Research Center comprise the current COSAM effort - "Conservation of Strategic

Aerospace Materials." The primary objective of COSAM is to help reduce the

dependence of the United States aerospace industry on strategic metals, such

as cobalt, columbium, tantalum, and chromium, by providing the materials

technology needed to minimize the strategic metal content of critical aerospace components with prime emphasis on components for gas turbine engines.

Thrusts in three technology areas are appropriate for COSAM. These include

near-term activities in the area of strategic element substitution; intermediate-range activities in the area of materials processing; and long-term,

high-risk activities in the area of "new classes" of high temperature

metallic materials. This paper describes in some detail the efforts currently

underway and initial results generated to date. Initial emphasis has been

placed in the area of strategic element substitution. Specifically, the role

of cobalt in nickel-base and cobalt-base superalloys vital to the aerospace

industry is being examined in great detail by means of cooperative universityindustry-government research efforts. Investigations are also underway in the area of "new classes" of alloys. Specifically, a study has been undertaken

to investigate the mechanical and physical properties of intermetallics that

will contain a minimum of the strategic metals. Current plans for COSAM also

are presented in this paper.

INTRODUCTION

The United States is heavily reliant upon foreign sources for the supply

of most strategic metals required by our aerospace industry. With the

exception of molybdenum, iron, magnesium, and the rare earths, the United

States imports from 50 to 100 percent of such aerospace metals Co, Cb, Ta , Cr

,

and. Mn (ref.l). However, the potential for foreign cartels, political unrest,

and production limitation is great and is intensified by steadily declining

known reserves. Thus, the United States can expect to be faced with supply

shortages and price escalation for many strategic metals. Since these metals

are vital to the welfare of the nation’s economy, their continued availability

at a reasonable cost is a national issue which requires cooperative action

between the aerospace industry and appropriate government agencies.

P4-1
The aerospace industry is currently a major factor in the positive inflow

of funds from U.S. exports (Ref. 2). This industry, and within it the aircraft engine industry in particular, relies heavily upon imports for several

key strategic metals including cobalt, columbium, tantalum, and chromium. In

order to offset or minimize future disruptions in supply, efforts. to develop

viable options must begin now, since a new material can take from 5 to 10

years of research and development efforts before qualifying for aerospace

service

.

NASA currently has plans to address the aerospace industry's needs to

minimize the use of strategic metals for advanced aerospace systems. COSAM

has as its broad objective the reduction of the dependence of the U.S. aero- space industry on strategic metals. This objective can be accomplished by

providing the materials technology options needed to allow individual com- panies to trade-off the material properties of critical components versus cost

and availability of their strategic metal content. This paper summarizes

NASA's current activities in this area and broadly outlines the plans for

COSAM.

STRATEGIC MATERIALS

A definition of strategic materials as used in this paper is given in

figure 1. Strategic materials are those predominantly or wholly imported

elements contained in the metallic alloys used in aerosoace components which

are essential to the strategic economical health of the U.S. aerospace industry. As the basis for what are considered strategic metals, we will focus on the aircraft engine industry's needs. Based on a survey of the ASME Gas

Turbine Panel and a subsequent survey of a number of aerospace companies, the

elements listed in figure 1 were considered to be the most strategic with

respect to the aerospace industry. As a result of prioritizing by NASA's

COSAM planning team supplemented by further discussions with several aircraft

engine manufacturers, four elements emerged that were of particular concern. The alloys used to build the critical high temperature components for aircraft

propulsion systems require the use of the four metals - cobalt, columbium,

tantalum, and chromium. These metals are contained in steels, stainless steels,

and superalloys that are used in engine manufacturing. Figure 2 lists these

four elements in the high priority category with a brief rationale for this

ranking. The remaining five strategic elements evolving from our surveys were

given a lower priority and figure 2 also contains a short explanation for this

ranking

.

The location of these metals in aircraft engine compressors, turbines,

and combustors is illustrated in figure 3. The need for such metals has increased as the demands have grown for higher durability plus high performance,

fuel efficient aircraft turbine engines. Based on the essential nature of

these metals and for the U.S. aircraft industry to maintain its competitive

position, it is necessary that supplies be readily available at a reasonably

stable cost. To achieve these requirements, domestic sources of key metals

are desirable.

P4-2
Today, we are almost totally dependent on foreign sources for these metals

as shown in figure 4. In several of the countries listed in figure 4, recent

political disturbances have led to supply interruptions. Therefore, the U.S.

aircraft engine industry can be seen to be highly vulnerable to supply instabilities of the essential metals for engine manufacturing. Accompanying supply

disruptions or increased demand is an accelerated price increase. Escalated

prices during the recent few years are evident for tantalum, columbium, cobalt,

and to a lesser degree for chromium, as shown in figure 5. These rapid price

increases illustrate the additional vulnerability of the U.S. aircraft engine

industry to cost fluctuations. The essential nature of cobalt, columbium,

tantalum, and chromium along with their vulnerability to supply instabilities

and cost fluctuations combine to cause these metals to be classified as strategic aerospace metals.

The portion of these four metals used in superalloys for the aerospace

industry compared to all other U.S. uses is shown in figures 6 through 9. The

use of these metals in superalloys as compared to total U.S. consumption in

1979 was: cobalt - 30 percent, columbium - 28 percent, tantalum - 5 percent,

and chromium - 3.4 percent. These data reveal that superalloys comprise the

largest single use of both cobalt and columbium.

OVERVIEW OF COS./U^l

COSAM has as its primary objective the reduction of the dependence of the

U.S. aerospace industry on strategic metals. COSAM can also provide the industry with some options for making their own property versus availability/cost

trade-offs when selecting aerospace alloys. These objectives will be achieved

by providing the technology needed to minimize the strategic metal content of

critical components in aerospace structures. Initial emphasis will be placed

on the aircraft engine industry. COSAM initially is focused on conservation

of the strategic metals cobalt, columbium, tantalum, and chromium. Strategic

metals such as titanium, the precious metals, tungsten, and others may be

brought into COSAM as it progresses.

Along with prioritizing the strategic elements that were identified, the

role that the NASA's COSAM effort should encompass was also evaluated. Options

that were considered are listed in figure 10. All of these options could con- tribute to the conservation of strategic materials and minimization of U.S.

aerospace industry vulnerability. However, within the scope of our program

a decision was reached based on LeRC traditional roles and expertise to focus

on the three areas noted in figure 10. These areas consist of strategic

element substitution, process technology, and alternate materials. Contributions to the other areas may benefit from COSAM through cooperative programs

with other governmental agencies such as in the area of scrap reclamation or through cooperation with technical societies in establishing a critical

material index. Having selected a list of four high priority strategic elements and having defined the areas of emphasis for COSAM and specific objectives, a technology approach was adopted as shown in figure 11. Conservation,

as well as reduced dependence on strategic metals, will be achieved in the

P4-3
area of strategic element substitution by systematically examining the effects

of replacing cobalt, columbium, and tantalum with less strategic elements in

current, high use engine alloys. This will help guide future material specifications if one or more of these metals becomes in short supply. Conservation

through process technology will be achieved by advancements in tl^ose net-shape

and tailored-structure processes that minimize strategic material input re- quirements. This will lower total usage. And in the longer term, development

of alternate materials that replace most strategic metals with those highly

available in the U.S. could lead to a substantial reduction in the U.S. dependence on foreign sources. Both of the later two technology areas will help

conserve the four strategic metals Co, Cb, Ta, and Cr.

EARLY COSAM ACTIVITIES

COSAM efforts began in FY'80. Efforts on planning and organizing are

still underway. In addition to the planning activities, several small re- search activities have been initiated. These research activities focus on two

of the three major thrusts of COSAM - strategic element substitution and development of alternate materials. Special emphasis of these initial efforts is on developing a fundamental understanding of the role of strategic elements in

current aircraft engine alloys so that effective alloying element substitution*

can be conducted. Similarly, in the development of alternate materials, a basic understanding of materials properties and alloying concepts is being

emphasized. Consequently, university grants play a major part in COSAM. In

addition, cooperative programs with industry augmented by in-house research at

the NASA Lewis Research Center comprise the approach used in these initial

projects. This cooperative approach will continue to be followed in COSAM

and industry, universities, and government in-house research will each play a key role. The subsequent paragraphs will describe in some detail early COS.A*'! research efforts.

Strategic Element Substitution

Four metals were mentioned previously as being classified as high

priority strategic metals. Cobalt was selected from these four metals for the

early COSAM strategic element substitution research. The basis for selecting

cobalt was twofold. First, the largest single use of cobalt in the U.S. is in

superalloys for jet engine applications as was shovsTi in figure 9 (Ref. 3).

Many of the other applications indicated in figure 9 are also important to the

nation's economy and security as well. Secondly, the specific roles that

cobalt plays in nickel-base superalloy fabrication and performance has not

been clearly established. Most superalloys currently in use were developed at

a time when cobalt was plentiful and inexpensive. Literature results (Ref. 4) are conflicting as to the role that cobalt plays in nickel-base superalloys in

important areas such as phase stability, y' partitioning, strength, fabricability, and oxidation and hot corrosion resistance. Because of these uncertainties, there exists a strong possibility that the strategic element cobalt can be

substantially reduced or possibly eliminated from several superalloys without

sacrifice of the key properties for which these alloys were selected for

engine service.

P4-4
^ Four nickel-base and one cobalt-base superalloys were selected for this

investigation. The five alloys are listed in figure 12 along with their

typical applications in the aircraft engine industry, the forms in which the

alloys are used, and remarks as to why they were selected for this activity.

Applications include turbine disks, turbine blades, and combustors. A variety

of product forms are represented by the applications of the five alloys as noted in figure 12. The selection of the five alloys was based primarily upon

the considerations given in this figure. Waspaloy* was selected because it

represents the highest tonnage of cobalt now in commercial aircraft engines.

Selection of Udimet-700* was based on the fact that this alloy is used in the

as-cast, as-wrought ingot, as-wrought powder, and as-HIP powder metallurgy

fabricated conditions. The potential for determining the impact of cobalt on both conventionally-cast as well as on single crystal turbine blades was the

reason for selecting Mar-M247*. Rene' 150* was chosen because it is one of

the most advanced directionally solidified alloys. The wrought, sheet alloy

HA-188* was selected because it represents one of the largest uses of a cobaltbase alloy in aircraft engines.

The primary purpose of the cobalt strategic element substitution research

is to determine the fundamental role of cobalt in a wide variety of nickelbase superalloys and in a high-use cobalt base superalloy. A secondary purpose

is to develop the methodology to explore the roles of other strategic elements

in similarly chosen alloys so as to have maximum impact on a wide range of

users

.

Figure 13 shows the participants in this COSAM effort on cobalt strategic

element substitution. These initial research efforts are planned for a threeyear period and consist of cooperative programs involving universities, industry, and NASA Lewis Research Center. Nominal compositions of the five alloys

given in figure 13 indicate that cobalt content ranges from 10 percent in MarM247 to 39 percent in HA-188. In addition, the y' phase ranges from 20 percent in Waspaloy to 65 percent in Rene' 150. The first phase in each research

effort will involve substituting the less strategic element, nickel, for

cobalt in incremental steps to a zero cobalt content. The effect of. this substitution on properties and phases present, such as y'> will make up the major

portion of the research effort in the first year of each program element.

Efforts in subsequent years will be directed at identifying and optimizing

alloying elements as substitutes for cobalt in the five alloys so as to maintain the key properties of these alloys.

The cooperative nature of the research being conducted on Waspaloy and

Udimet-700 is illustrated in figure 14. The role of industry as represented

by Special Metals Corporation is outlined. Their primary role is to characterize and optimize fabrication and heat treating procedures for the reduced

*Trademarks

Waspaloy

Udimet

Mar-M

Rene

HA

United Technologies Corporation

Special Metals Corporation

Martin Marietta Corporation

General Electric Corporation

Cabot Corporation

P4-5
cobalt Waspaloy and Udimet-700 alloys. The university role in this effort is

also shovm in figure 14. Columbia Uniyersity will be involved with mechanical

property characterization, structural stability, microstructural features, and

theoretical formulations to identify future alloy modifications if required

for the second phase of the project. Purdue University will be pr-imarily

responsible for microstructural and microchemistry characterization of the

reduced cobalt content alloys. To round out the program, NASA Lewis Research

Center will be involved in further mechanical and physical metallurgy characterization of the alloys as shown in figure 14. The output of this cooperative effort is expected to be a clearer understanding of the role of cobalt in

nickel-base superalloys.

Some preliminary results on the effects of reducing cobalt in Waspaloy,

a 13 percent cobalt alloy, are shown in figure 15 (Ref. 5). Tensile strength

appears to be insensitive to the amounts of cobalt in the alloy. However,

rupture life decreased with decreasing amount of cobalt in Waspaloy. Further

testing will be required to better characterize this apparent effect.

Similar effects of cobalt on the rupture life of Mar-M247 have been deter

mined as shown in figure 16 (Ref. 6) . A possible contributing cause to this

reduction in rupture life is the decrease in amount of Y’ in this alloy with

decreasing cobalt content as shown in figure 17. Also shown in figure 17 is

the change in y' composition. As cobalt is removed from the alloy, the

largest change in the composition of y' is the increase in tungsten content.

Further studies are underway to clarify the role of cobalt in this alloy.

The research efforts on Udimet-700 and Rene' 150 parallel the previously

described efforts on Waspaloy and Mar-M247. It is anticipated that these

studies will lead to an understanding of the fundamental role of cobalt in a variety of conventional and directional nickel-base superalloys. These

results should provide an improved technical base to develop modified superalloys in later stages of COSAM, as illustrated in figure 18.

Alternate Materials

Research in this area must be considered to be high risk and long range,

but it has the potential of a high payoff in terms of significantly reducing

the nation's dependence on strategic materials. As an example of alternate

materials, intermetallic compounds are currently being investigated for

possible structural applications. Initial efforts are centered on nickel and

iron aluminides. Successful development of this type of alternate material

offers the possibility of partially or totally replacing all the strategic

materials in components where intermetallic compounds can be utilized.

Intermetallic compounds are of interest because of their potential high

temperature strength as shown in figure 19 (Ref. 7). It can be seen in this

figure that nickel aluminides have the strength capability of competing with

current nickel-base alloys. However, a possible disadvantage of this type of

material is that simple binary aluminide compounds have shown a lack of room temperature ductility (Fig. 20). The factors which influence the high ductile

P4-6
to-brittle transition temperature of nickel alurainide C~600 + are currently being investigated. A NASA grant with Dartmouth University is

aimed at understanding the fundamental deformation mechanisms in nickel

aluminide. From these investigations, methods of improving the low temperature ductility of nickel aluminide may be suggested. An accompanying inhouse research project at NASA Lewis Research Center is focusing on the high

temperature mechanical properties of aluminides. These studies can provide a fundamental basis for more extensive research to develop these nonstrategic

,

alternate materials as shown in Figure 21.

COSAM PLANS

Future COSAM efforts can build on the fundamental understanding from the

early research for cobalt substitution, as was shown in figure 18. Major

efforts will be devoted to developing, and if warranted, to scaling-up low or no-cobalt nickel base superalloys for fabrication into various components.

Demonstration of continued promise could also lead to verification in engine

tests by major engine producers. Similar efforts will also be conducted for

other strategic metals such as columbium and tantalum.

In the area of alternate materials, much more work will be required to

develop materials such as intermetal lie compounds. As was shown in figure 21,

initial efforts will focus on fundamental studies aimed at improving low

temperature ducitility and high temperature strength of FeAl and NiAl intermetallics. Complete property characterization will follow on more promising

compositions. Reiterations of these basic steps will be required to further

optimize the alternate materials and make them viable candidates as structural

materials for aircraft engines. Scale-up and rig testing of promising compositions for blades and vanes will follow. The development of alternate

materials will help conserv'e the strategic metals Co, Cb, Ta, and Cr.

A third area of the COSAM consideration involves conservation through

improved materials processing technology. Although none of these activities

have been initiated, plans have been made for investigating processing

technology in such areas as advanced melting techniques, tailored fabrication,

advanced coatings, joining techniques, and fabrication efficient processes. A

reduction in strategic material usage should result from these processing

technologies. For example, early efforts on near-net-shape fabrication of a turbine disk (Ref. 8) have been shown to be able to reduce input material

weight compared to conventional casting/forging practice and further gains

appear possible. Improved processing technology will also help conserve the

strategic metals Co, Cb, Ta, and Cr.

CONCLUDING REMARKS

This paper has presented NASA's COSAM efforts and planning. The primary

points are summarized below:

P4-7
1. Advancements in materials technologies are needed to provide the aerospace

industry with alternative materials options in the event of future strategic

metal shortages or excessive price increases.

2. The primary role of NASA's COSAM efforts will be to address strategic

material problems within the aerospace industry. COSAM should make contributions to a national data base that will benefit many other domestic industries

as well.

3. COSAM was designed to involve cooperative research efforts with industry

(alloy producers, component fabricators, and engine manufacturers), with

universities, and with government research facilities (primarily the Lewis

Research Center).

REFERENCES

1. Mineral Commodity Summaries 1979. Bureau of Mines, U.S. Department of

the Interior, 1979.

2. "Aerospace Review and Forecast 1979/80", Aerospace, Vol. 18, No. 1, Winter,

1980, p. 4.

3. Mineral Industry Surveys, Bureau of Mines, U.S. Department of the Interior,

December 1979.

4. Tien, J, K. ; Howson, T. E.; Chen, G. L. ; and Xie, X. S. : "Cobalt Availability and Superalloys", Journal of Metals, October 1980, pp. 12-20.

5. Maurer, G. E.; Jackman, L. A.; and Domingue, J. A.: "Role of Cobalt in

Waspaloy", Superalloys 1980, J. K. Tien, ed. , ASM, Metals Park, OH,

1980, pp. 43-52.

6. Nathal, M. : "The Role of Cobalt in a Nickel Base Superalloy". Presented

at the AIME Annual Meeting, February 22, 1981, Chicago, IL.

7. Schulson, E. M. ; "The Ductility of Polycrystalline NiAl", Presented at

NASA Lewis Research Center, September 1979.

8. Dreshfield, R. L. ; and Miner, R. V., Jr.; "Applications of Superalloy

Powder Metallurgy for Aircraft Engines", 1980 International Powder

Metallurgy Conference, Washington, DC, June 22-27, 1980.

P4-8
COSAM

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P4-12

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P4-13

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CONSUMPTION

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P4-15
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P4-16

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PS-SO-M72
CHARGE TO WORKSHOP IN RELATION MATERIALS

POLICY LEGISLATION (PL 96-A79)

John B, Wachtman

National Bureau of Standards

4

P5
Workshop on Conservation and Substitution Technology

for Critical Materials

Vanderbilt University

June 17, 1981

John B. Wachtman, Jr.

Director, Center for Materials Science

National Bureau of Standards

Charge to Workshop

This workshop is being held primarily to assist in the preparation

of a report to Congress required by a new law passed last October, the

National Materials and Minerals Policy, Research and Development Act of

1980. This Act gives assignments for separate reports to the Departments

of Commerce (DOC) and Defense (DOD) both due for delivery to Congress on October 1, 1981. It also assigns other responsibilities to the Department

of Interior and to the President's Office of Science and Technology.

Some ten agencies are assisting us with this DOC report, indeed the

Department of Interior is a cosponsor of the present workshop. The

present workshop is being held especially to develop information for the

report required by the Department of Commerce, but should be very useful

to the other agencies in their responsibilities also.

Slide 1 (P5-4). To give a broad perspective on the present workshop,

let's take an overall look at the two reports which the Act requires.

The DOD report is supposed to address critical material needs relative

to defense and to identify items to meet their needs.

The DOC report is supposed to identify a specific materials

needs case related to national security, economic well-being, and industrial

productivity, to assess critical materials needs and to recommend programs

to meet these needs.

Slide 2 (P5-5). The DOC report is narrower than the DOD report in one area and broader in another. It is narrower in the sense of being a case

study. We selected the aerospace industry for the case study. With the aid

of industry representatives and other experts, we held a workshop in

Gaithersburg, Maryland on February 9-10, and selected six specific materials

to typify the range of problem and opportunities associated with the critical

materials problem.

The DOC report is broader than the DOD report in being

explicitly concerned with industrial productivity and economic well-being

as well as national security.

Slide 3 (P5-6). The DOC report will analyze the issues associated with

supplies of all these materials including supply in all forms and stages

of use; i.e., including parts and assemblies, not just raw materials.

We want to know what policy options are available to the Government and

which ones it should pursue. That is, what should the Government do or stop doing or what combination of these.

P5-1
Slide 4 (P5-7). We clearly recognize that there are a whole spectrum of

ways to respond to a supply problem, depending on the nature and severity

of the problem. One set of supply oriented options would include increased

domestic mining, increased stockpiling of materials, and a foreign

policy directed more strongly toward security of supplies. These options

are under study by the Department of Commerce team but are not the focus

of the present workshop. We are concerned here with the second class of

options; the technical options. These include substitution which we consider in two categories. First, substitution which taken place

naturally as technology advances and the economics of new materials and

processes leads to their use. We are calling this economically favorable

displacement or just displacement. The second kind of substitution

takes place in an emergency and may cause higher cost or lowered but

adequate performance. We are calling this emergency substitution or sometimes just substitution for short.

Our present workshop is designed to get at opportunities

and recommendations in this area of displacement, substitution, and

conversion. We do not want to spend any time at this meeting on discussion

of mining, stockpiling of materials, and the like because they are being

treated by other parts of the Commerce team. We do want to get your

ideas on where the most promising technical opportunities are for reducing

our vulnerability to problem associated with critical imported materials,

especially through the development of advanced materials. Bob Nash has

prepared some questions which we hope you will help us with in the

so-called "brain writing" sessions.

Slide 5 (P5-8). Many other agencies are assisting the Department of

Commerce in preparing its report. We have asked two of these, the

Department of Defense and the Department of Interior, which are named in

the Materials Act of 1980 in connection with the Department of Commerce

responsibility, to give brief summaries of their Research and Development

program relating to critical materials. We have also asked for a presentation from the National Aeronautics and Space Administration because of

its obvious importance to the Aerospace Industry. The other agencies

listed on this slide have representatives here also in most cases. We

appreciate their help and look forward to their active participation in

the Workshop.

P5-2
PL 96-^79

DOD Report:

1) Assess critical mTERiALS needs related to

flATIONAL SECURm.

2) Identify steps to meet needs.

DOC Report:

1) Identify specific materials needs case reu\ted to

NATIONAL SECURITY^ ECONOMIC WELL-BEING^ AND

INDUSTRIAL PRODUCTION,

2) Assess critical materials needs in that case and

RECO^T1END PROGRAMS.

3) Continue to assess additional cases.

P5-3
DOC REPORT

0 Narrov/er in focus: a specific case — We have selected the aerospace industry

-- Within aerospace^ as examples:

1) Chromium 4) Tantalum

2) Cobalt 5) , Rapid solidification technology

3) Titanium 6) Composites

0 Broader in another sense — Economic — Industrial^ as well as — Security

P5-4
DOC REroRT

0 Issues associated with supplies of

SIX mTERIALS

0 Policy options

0 Program REca^wENDATioNS

P5-5
CLOSE COORDINATION

1) Federal emergency management agency

2) Department of the interior

3) Department of defense

A) Central intelligence agency

5) National aeronautics and space administration

6) Department of Energy

7) Department of state

8) Department of transportation

9) National science foundation

P5-6
WORKSHOP GOALS

fV\iN Goals;

Subsidiary Goal:

Define technical opportunities for — Dispu\cement — Conservation “ Substitlition

OF IMPORTED MATERIALS

Identify targets for "Information Stockpile"

Convey views on federal action that would

BE DESIRABLE TO FACILITATE THE DEVELOPMENT

OF ADVANCED MATERIALS.

P5-7
CONSERVING CHROfllUri

David C. Goldberg

Westinghouse Electric Corporation

P6
KEYNOTE ADDRESS:

June 1981

CONSERVING CHROMIUM

D. C. Goldberg

Westinghouse Electric Corp.

In preparing this overview talk I have drawn heavily on the NMAB-NAS

study "Contingency Plans for Chromium Utilization" issued in 1978 as

the data base, the summary, and the recommendations contained in it

are as valid today as it was at that time.

Another excellent resource was the book "World Mineral Trends and

U.S. Supply Problems," Resources for the Future, R-20, December 1980.

The authors of this book highlighted the mineral chromium as one having

an undue vulnerability to supply disruptions and subsequent economic

impact.

It is a well known fact that, with the exception of Russia, no developed

country, which has a large market for materials dependent on chromium,

has large deposits of chromite. These countries have had to import such

large quantities of chromite that the supply-demand scenario can, at

various times, reach "crisis" situations. Given the current trend of

increasing U.S. dependency on the continuing availability of South African

chrome ores, it is significant that the current administration is

fostering an environment, both through its actions and deeds, that is

considerably more favorable for developing a strategic minerals policy

that will go beyond the long prevailing and inadequate "emergency" concept.

We also see the United States Senate and House moving forward on implementing the provisions of National Materials and Minerals Policy,

P6-1
Research and Development Act of 1981 (P.L. 96-479) signed by President

Carter on October 21, 1980.

The long range commitment to a strategic minerals policy was emphasized

when Secretary of Interior James Watt recently told the House Mines

and Mining Subcommittee (March 26, 1981) that a strategic minerals

policy is "not simply the establishment of emergency government action"

but "policies that protect American jobs and investments, improve our

balance of trade, revitalize the Nation's economy, and provide for the

security of foreign mineral imports." He also told the Senate Energy

and Minerals Subcommittee the following week that "it is imperative to

take a national strategic view of the land use in the focus of mineral

requirements 20 years or more down the road."

These statements plus administration actions* are aimed at ending South

Africa's "polecat status" in the world if they cooperate in achieving

an international ly-acceptabl e settlement for the independence of

Southwest Africa. All of these moves provide a climate which could

diminsh the proportions of our current dependence on mineral supplies

from hostile or politically sensitive areas.

It is fitting that Chromium is the first of the metals to be discussed

in this three-day workshop as it is probably the one metallic element

for which there is not an obvious substitute. The papers to follow,

review recent actions in conservation and substitution.

*New York Times News Service - 5/30/81

P6-2
Before we get into these specifics, I would like to present an overview of recent studies on this important element. Chromium is an

element that imparts unique properties to the material (s) to which

it is added, but it is not scarce and its price has been relatively

low until recently. It plays a vital and almost irreplaceable role

in widely disparate commodities, and chromium containing materials

are used in a variety of essential military and industrial applications.

I thought it would be a good idea to begin the discussion on an historical note to remind us that the problem we are facing today is not all

that new. Certainly, many of the rules of the game have changed.

Energy costs and environmental factors have a lot to say about the

materials we use today to manufacture goods.

But the basic solutions to the problem — identifying, substituting,

recyclying, stockpiling, exploring for alternate materials and developing

new processes to refine and manufacture materials — have remained

pretty much the same.

The United States faces only one type of "major mineral supply problem":

an undue vulnerability, of some minerals due to contingencies that might

either seriously disrupt supplies or cause a sharp upward movement of

prices, with serious economic impacts. This is particularly true for

chromium, in the form in which is is consumed by the steel industry

(principally ferrochromium)

.

P6-3
It is projected that United States consumption of chromite will decline,

although at a lesser rate than in the recent past, while world consumption

will increase by about 70 percent. United States ferrochromium consumption is expected to increase by approximately 60 percent, more than

double in Japan, and quadruple in the Soviet Union. Accordingly, we

will see some of the same economic and political measures used in a

•self-serving manner, to insure their uninterrupted ore supplies.

Chromium also presents no basic resource problem, and it appears as if

the present tightness in mining capacity will diminish over the balance

of the century. There may be some difficulty, however, in maintaining

an adequate supply of refractory-type ore, owing to concentration in the

Philippines. Ferrochromium is likely to become increasingly concentrated

in southern Africa.

Chromium exemplifies the problem of long-term exposure to supply contingencies, which occurs when there is physical concentration of supply

in areas (South Africa) that are unusually subject to politically motivated supply disruptions. Higher energy costs and environmental problems

have driven United States ferrochromium processing to South Africa also,

thus adding to our vulnerability to a supply interruption.

Chromium is also exposed to short term contingencies of embargoes and

racial conflicts in the countries of chromium origin. The ample supply

of chromite ore offsets the impact but is still a potential problem

with respect to the ferrochromium. Given the global political situation.

P6-4
the possibility of forming a chromium cartel that could produce severe

economic penalties and industrial disruptions by interrupting United

States chromium imports is of serious concern.

Although chromium is found in various minerals, the ore known as

chromite is the sole commercial source. Historically, chromite has

been classified into three general categories; metallurgical, chemical,

and refractory. Technological advances permitted considerable interchangeabi 1 ity among the various grades, particularly in respect to the

chemical grade, which can be used in all industries. The principal

substitution has been to chemical for metal lurgical . Refractory (high

alumina) ore has tended to dominate its market.

Because of complacency, materials research has not received the proper

level of attention and funding it deserves. Increased awareness of

materials during periods of war resulted in increased fundi ng of

research. Out of this effort, of course, came synthetic rubber, new

electrical insulating materials, a host of new alloys, polyester

fibers, polyethelene, revolutionary molding compounds and other

materials.

In many areas of materials research, we are going to have to run very

fast just to stand still — just to hold on to the quality of life we

enjoy today. This is particularly true in the metals area.

P6-5
I think the future of materials research work will be devoted to the

development and refinement of alternate materials, and to the exploration of ways in which we can make more effective use of the materials

we do have.

Expenditures in materials research and development in this country

peaked in the 1960's. While current dollar annual expenditures in the

field are five times what they were in the 1950's when viewed from the

perspective of inflation, investment in materials research has actually

been on the decline for several years. As is the case with so many

other areas of research, we have definitely seen a decline in innovative

risk taking in the past 10 years. As a result, R&D in the field has

been much more selective and cautious.

It is true that when we realize we have a severe shortage, R&D has

been funded. Cobalt shortages, for example, have provided the economic

incentives for industry to seek substitutions. The electrical industry

uses cobalt to make magnets, but we are learning to make rather remarkable

magnets without cobalt.

When R&D has not been funded, we are in trouble. Chromium is perhaps

the classic example. We as a nation are behind the times in chromium

substitution research. This is largely because of the research climate

that exists today in this country — a climate that is burdensome to

industry, and quite different from the cooperative climate that exists

between industry and government in countries such as Japan and West

Germany.

P6-6
Government policy has been channeling resources into energy research,

which is obviously vital to our well-being. But government, industry

and the universities have all virtually ignored materials research.

It is clear that in order to ward off an impending materials crisis,

we must make materials research a national priority. Materials research

has been largely directed to knowing more about the materials we have.

But little has been done to search for replacements and to explore

totally innovative ideas.

Our government has been for years shooting itself in the foot insofar

as the health of our domestic mineral industries is concerned. In

order to stimulate the necessary domestic supply of critical materials,

the government will have to get its act together in the form of a

national materials policy.

For too long uncoordinated government policies on land use, environment,

stockpiling, taxes, and depreciation allowances have been made without

weighing their adverse effect on domestic mining and processing of

critical materials. We now see ferro-alloy plants shutting down in the

United States and opening in South Africa, and a copper refinery shut

down here and the operation moved to Japan, making us more rather than

less dependent on foreign sources. Even our foreign policy has at times

jeopardized the flow of critical materials. All of these policies must

be changed so that full recognition of critical materials concerns is

given, hence the need for an overall national materials policy that

becomes a factor in our foreign policy negotiations.

P6-7
The same rationale, I believe, applies to materials. The only way to

prevent the formation of an OMEC is to develop new materials, and new

materials processes, that will decrease the bargaining power of the

sub-Saharan nations.

Like the energy problem the materials problem ultimately boils down to

a basic problem of supply and demand. It is fundamentally a problem

of economics that is addressed in both short and long term planning of

business and government working in partnership.

And since when it comes to materials, the United States is a have-not

nation, ingenuity is the only card we have to play.

Our mineral resources, after all, are finite on our planet. Our

challenge today is to get smarter about the minerals and materials we

do have. We have to learn how to extract these minerals more efficiently

from nature. We have to learn how to produce materials from these

minerals using less energy. And we have to learn how to recover more

of the materials we use in society.

These are all challenges for technology. Government policies can help.

Certainly economic and energy factors will play key roles in our future

material consumption habits.

But technology remains the driving force required to transform the

finite resources of our planet into economic realities.

P6-8
The efforts of technology, however, must be supported by prudent

government planning and increased public awareness.

From the conclusions of the various authoritative studies made on

chromium in recent years, it is clearly apparent that a supply disruption of several years would create serious — both short and long

term — economic problems. When this is viewed in terms of where the

world's chromite supply will be located in the next century, I feel it

is necessary to initiate programs with respect to exploration, stockpiling, recycling, substitution, and alloy development at an accelerated

rate.

P6-9
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P6-22
COMPUTER SYSTEMS FOR DEVELOPING CHROMIUM

FREE STEELS FOR HARDENABILITY

Dale H. Breen

International Harvester Co.

P7

COMPUTER SYSTEMS FOR DEVELOPING CHROMIUM

FREE STEELS FOR HARDENABILITY

Dale H. Breen

International Harvester Co.

P7-1
Good morning, Gentlemen:

It's a pleasure for me to be a participant in this work shop and conference concerning strategic alloys. The subject that I am going to

discuss, substitution based on hardenability , is indeed an appropriate

one to be in this conference although it addresses only a small segment

of the problem. I will have a number of exhibits (attached) that I'm

going to use to help me demonstrate to you a cost motivated technology

that we have used at the International Harvester Company for some time.

The technology has broader inferences than that, however, and that is the

purpose of my being here, to explain how it is applicable to threatened

shortages of strategical alloys.

Fig. 1 shows a condition that has already been well-defined in the

literature so I'll not dwell in this area. It demonstrates the location

of our chrome resources and shows, in this instance, that the Western

Hem^isphere doesn't have very much. As can be seen, the seven million

one-hundred thirty thousand short tons represents only a very small

fraction of the world total. Fig. 2 summarizes the condition of the

Eastern Hemisphere and the countries that are in that portion of the

world. As you can see, their hemispheric total reserve of a billion

eight hundred and thirty-four million three hundred and eighty thousand

tons is quite a high fraction of the world total. Also one can see that

it is mainly concentrated in South Africa and Southern Rhodesia but

Russia also has a significant amount. It is not my purpose here to get

into the political aspects of this but the very least one can say is that

there is potentially a problem in obtaining chromium.

P7-2
- 2 -

Now to move into the area that I want to address, that is, in the

application of chrome-bearing steels for ground vehicles. There are

many, many tons of steels used in structural applications, such as gears

and shafts. Referring to Fig. 3, they are designated as chrome, chrome

vanadium, nickel-chrome, nickel-chrome-moly-vanadium, chrome-moly and

so forth as shown. These carry SAE or AISI numbers, such as the straight

chrome steels which are the 5100 series, the nickel-chrome-raoly steels

are 4000 and 8600 type steels and so forth. The object of Fig. 3,

though, is to show how many tons of chromium are required to make these

steels based on 1978 production. This tonnage represents somewhere in

the order of 10 to 12^ of the total chromium used domestically. None of

these steels are heat resistant or stainless steels. These are steels

primarily for use in structural applications such as gears and shafts

(refer Fig. 4). The technology associated with the engineering of such

components is quite developed and very sophisticated. Fatigue strength,

wear resistance and dimensional control are very important. They are

applied in ground vehicles of all sorts. Automobiles, of course,

highway trucks, earth-moving equipment, locomotives, etc. all use these

type of steels in relatively high tonnage.

Now, another important factor is that the single most important aspect in

their applications is their hardenability characteristics. They are

generally purchased to standard hardenability bands rather than

chemistry. To explain further, the point I want to make now is that

worldwide, steels are bought and sold to hardenability bands (refer

Fig. 5) and that hardenability is ^ne of the most significant criteria

P7-3
- 3 -

related to their performance in applications. Hardenability is the ease

with v/hich they may be hardened. This will be defined more explicitly

latei' but the point 1 want to make is that tiiere are standards, and that

the standards are defined in terms of hardenability. Now I’d like to

dwell in the area of semantics a minute because I am going to go in two

different directions. When I talk substitution technology, I talk about

tv7o different concepts; replacement steels and alternate steels (refer

Fig. 6). A replacement steel is a steel, by my definition, that has all

of the essential hardenability characteristics of the steel it is going

to replace. In other words, the hardenability bands are identical. The

chemistry will be different because it's a different steel but the

hardenability bands will be essentially the same. An alternate steel on

the other hand, may be successful in a given application, but it will be

a different hardenability steel. The hardenability that it would have,

however, would be adequate for the application. I'm going to remind you

of this difference as we proceed and give specific examples.

First though, I am going to talk in terms of substituion in general. I'd

like to give a little bit of background on where we've been historically

on substitution and a little bit of background about hardenability. This

portion won't be "in-depth", just enough to allow us to follow through

some of the subsequent reasoning. Fig. 7 is a very significant graph

related to substitution. It shows the usage of the alloy steels types

that we are considering over a period of seventy years. You can see

here starting back in the 1900 era that the straight nickel steels were

P7-4
- 4 -

very popular. They gave way in popularity to the chrome-vanadium steels,

and the nickel-chrome steels. A very significant point on this graph is

the area where the nickel-chrome-moly steels became popular. These are

8600 steels that have been the standard of the industry in the types of

application we are discussing since the early 19^0 's. These Ni-Cr-Mo

steels are really substitution steels. They were at that time referred

to as national emergency, NE, steels. Nickel was in short supply and

nickel-bearing steels simply had to be replaced. So this shows we have

been sucessful in this endeavor of making substitutions before. It was

quite painful though. The literature of the era says that it was done

with a great deal of fear and trepidation. In addition, the implication

of this graph is that many of the changes that have occurred have been

promoted by catastrophes (e.g.) World War I, World War II and the Korean

emergency. All caused changes in the approach to using steels due to a

lack of access to certain alloying constituents. We at International

Harvester Company have studied this system, seen that it works when it

has to, and decided that it could work when least cost was the principal

driving force. So we have developed a systematic program which I am

going to describe for you that addresses the issue of substitution where

least cost is the goal. But before we leave this chart (Fig. 7) I'd

like to quickly talk in terms the horizontal axes, the time line, and

some of the technologies that allow us to make the substitutions that

v/ill be described subsequently; i.e., the technology base. For

instance, in the mid 1925's Harry McQuaid did his famous work on grain

size. He established a grain size measuring method and demonstrated the

P7-5
- 5 -

effect of grain size on hardenability . In the early 1930’s, Edgar Bain

began to unravel some of the intricacies of transformations and we

learned a lot more about hardening rationale. In the late 1930's, Walter

Jominy developed the end-quench test which is the major tool in use for

measuring hardenability. It has been used since that time on a

worldwide basis as the criteria for buying and selling steels. During

this period also Marcus Grossman developed the basis of our quantitative

substitution system. He developed m.ultiplying factors for each

individual alloy to enable the calculation of hardenability. Subsequent

to that, in the late 40 's there was a considerable amount of work

building on Grossman's earlier work. For instance, Kramer, Segal and

Brooks and others published additional quantitative data. Then there

was a dry spell until the late 1950's when Girardi and Jatzak provide

multiplying factors for individual elements for high carbon steels and

carburizing steels. They pointed out that alloy effects were different

at low carbon levels than at high carbon levels for the same alloying

element. For instance, if one were trying to calculate expected

response in the case portion of a carburized part, he would be very much

in error if he only considered base material hardenability factors. Then

there wasn't much done till the late 1960's when Doane did some more

work on multiplying factors. A considerable number of people, Kirkaldy

and others, and ourselves, have developed regression equations to

calculate hardenability and then, in addition, our own work in applying a

separable computer program to design least cost alloys. This then

reviews the background and the way this technology has developed over a

long period of time.

P7-6
- 6 -

I'd like to, before I go any further, bring out one of the major problems

related to substitution. To bring it into focus, refer to Fig. 8. I

picked up this headline, "Pass the Aspirin: We're changing materials,"

out of a trade magazine, but it reflects the general attitude in the

industry, i.e., the engineering and manufacturing community. Once a

material has been proven, the engineer likes to stick with it. He

doesn't like to make changes. He will, if that material performs

exceptionally well and it develops a warm spot in his heart, even use it

in places where it shouldn't be used. His concern is legitimate. He

has something that works. The design art is so complex that he just

doesn't like to touch things once they are working well. In these days

of consumerism, the possibility of product litigation and call backs

it's proper to proceed with caution. Handling this attitude will be

discussed later.

I would like now to describe a case history of our use of this technology

and I am sure, through this experience, you will be able to see how it is

applicable to the chromium situation. The system we use is based on the

concept that if hardenability remains the same, the same microstructure

will result on heat treatment and performance characteristics will be

equal. Our system is called the CHAT system (refer Fig. 9), an acronym

meaning Computer Harmonized Applications Tailoring, and it is a

quantitative system for the design and selection of heat treatable alloy

steels. It has two parts. The AT part. Application Tailoring, which I

am going to discuss last but I'd like to define first (refer Fig. 10 ), is

a formalized approach used to dete. mine the hardenability levels

P7-7
- 7 -

required to develop strength gradients required for given applications.

That is the system we would use to determine the metallurgical

requirements for a given design. The CH portion (Fig. 11), which will be'

emphasized, is the computer selection of the least cost steel analysis

which will meet the m.inimum hardenability level called for by

Application Tailoring. If a part has already been Application Tailored

and the hardenability band requirements are set, then one can use the CH

process to develop the least cost replacement composition that will have

those hardenability characteristics.

Let’s review quickly some elementary aspects of hardenability. This is

an end-quench bar (Fig. 12). It is heated then quenched on one end to

measure hardenability. The dark end is being quenched. So one end is

fast cooled and the other slow cooled. After it’s completely cooled to

room temperature, flats are ground in the longitudinal direction and

hardness indentations are made along the length of the bar. The hardness

readings are then plotted against distance as is demonstrated in Fig. 13.

I want you to keep in mind that many parts made from the steel grades are

machined out of low carbon steel and then placed into carburizing

furnaces, carbon diffused into the surface and then quenched so that they

are in effect two different steels; high carbon on the surface and low

carbon in the core. Near the quenched end the steel is harder and then

becomes softer towards the slow cooled end. Now, Grossman’s work said

that this curve in Fig. 13 could be predicted on the basis of multiplying

factors for the alloying elements that were involved in the particular

piece of steel. He calculated what’s called a Dj. We’ll not explore Dj

P7-8
in depth. It’s defined as the ideal critical diameter but that’s not

relevant for our purpose. Let’s think of it in terms of an index number

f for hardenability. Anyway, the Dj is a function of carbon and alloy

content and can be calculated using multiplying factors. Referring

still to Fig. 13 , initial hardness is a function of carbon and the

distance hardness is a function of carbon and alloy content. The

rapidity of drop off relates to Dj. High Dj’s, high alloy contents are

characterized by small changes in hardness with distance and low. Dj’s

drop off closer in as is shown. Fig. 14 shows two different groups of

steels. They’re grouped by base carbon level. The higher carbon level

will start with a higher surface hardness. A low DI will drop off faster

and a high DI will be harder farther out on the bar. A lower carbon steel

would start at a lower hardness but could have similar Dj . So its

hardenability could be the same, i.e., harden to the same distance but

have a lower initial hardness and lower distance hardness due to being

lower in carbon. The concept I am establishing is that the hardenability

of steel can be calculated, not only by this method, but by other methods

as well.

Now what we would like to point out though is, that steels are not simple

chemistries, but are ranges. They’re made day after day as 140 to 240

ton heats and chemistries and hardenabilities vary from heat to heat and

within the heat. Mills have agreed, however, on the basis of experience

to meet hardenability band specifications. This particular band (refer

Fig. 15) is labeled EX 6 but it’s the same band as 8622H. EX 6 was one

of the early replacement steels we developed and it has exactly the same

P7-9
- 9 -

hardenability characteristics as 8622 so I am going to use it for my

example. Now you can see, taking the mid-range chemistry, that the

calculated hardenability splits the band pretty well. But when one takes

the chemistry extremes that are allowed, we’re inaccurate on both sides

of the band. This, of course, relates to the manner in which

hardenability bands are established. The principal reason is that the

chemistry is seldom all low side or all high side in practice; hence,

it’s really a statistical problem. We’ve looked into this area by

examining the statistical distribution (refer Fig. 16) of individual

alloy elements in a number of heats. We decided to try to calculate the

range based on some variation in sigma in this case +2<^. We recognize

though that the variation in +2

^ chemistry may give different range of

m.ultiplying factors depending on v^;hat percent of element is involved

(refer Fig. 17). I mention this factor as a matter of information. But

when we applied +2 sigma concepts, we were able to do a fairly reasonable

job of calculating the hardenability band for the 8622 and Ex 6 (refer

Fig. 18). Now I’d like to point out something here, the Djg for midrange is 2 . 3 . The Djg, +2 ^, is then 1.6 to 3.1* As yet, we haven’t

considered the Djq . The Djq is the hardenability at high carbon

levels, say .90 carbon. This low carbon steel {. 22 %) is going to be

carburized, so we have to know what the hardenability is at high carbon

and Djq's have to be matched. You might keep that in mind, i.e., there’s

another curve set like this that represents the Dj in the high carbon

region. Just keep in mind the ^4.8 from Fig. I 8 because it is involved in

my example of a replacement steel which was developed during the nickel

shortage in 1969. Remember a repl;. cement steel has exactly the same

P7-10
- 10 -

hardenability characteristics, both case, Djq, and core. Dig. Our

objective is to use the principles of CHAT to design a least cost steel

with no nickel. I want to describe the CHAT concept in more detail

first, however. In CHAT, of course, one of the basic concepts is relative

hardenability efficiency, defined as the log of multiplying factor

divided by cost. It's simply how much kick in hardenability is

involved and how much does it cost. Figs. 19 and 20 are some curves that

give relative hardenability efficiencies of some common alloy elements

used. You see the manganese is very attractive in terms of low carbon

base hardenability, chromium is also, moly and nickel aren't as

efficient. If these data were updated with recent changes in price of

moly, moly would not look even this attractive. In the high carbon

region as you can observe silicon is m.uch more attractive, chrome is even

better than manganese at certain levels of alloying element, etc. This

is the type of information that's in our computer program. We are in

effect saying to the computer, here are the hardenability requirements,

give us the chemistry that would provide these hardenabilities , both

case and core, at least cost. In this example. Fig. 21 , we did just

that. At the same time, we restricted Ni. The composition of the new

low Ni steel is also shown. It' a little lower chrome, higher moly,

significantly higher manganese, but you notice that the Djg is

essentially the same, and the Djq, at 4.7 is essentially the same. So

you v;ould expect this alloy steel to perform just as well in the

applications we had in mind for as the other steel. But the engineering

community doesn't agree with this sort of thing. If you assign another

number to a steel, to them that's different steel even though you m.ade

P7-11
- 1

1

only minor changes. That's where the "Pass the Aspirin" concept comes

in. It's really a fixed concept in the minds of most engineers. So we

have to do some test work to provide data to show equivalency. I'm

going to review for you quickly the results of some test work on this

steel. Ex 16, by the way, is also called Equalloy. The EX designation

comes from SAE. Before I go into that though, I'd like to show where one

can bo led asti^ay. We have been talking about Djp and Djg and I told you

that Ex 15, Ex 16 type were the same in both case and core hardenability

as the steels they replace. I showed you the core hardenability

previously, but now here's the results of a test in the high carbon region on this steel compared to 8622H (refer Fig. 22). This demonstrates

the similarity in case hardenability to 8622. The other alloy here, Ex

10, has identical base hardenability characteristics as 8622 but

significantly different hardenabilities in the high carbon regions. Now

this would be a disaster if the requirements on a gear, for instance, was

a cooling equivalent of 12/16" on an end-quench test. Only 50 Rockwell

would be achieved on the surface in the carburized regions with this

steel whereas you would be getting the high hardness you'd expect in Ex

16, giving the high fatigue and wear resistance characteristics, needed

for gears. So one has to recognize £hat both of these characteristics,

Djg and Djq, are important. Our motivation as I said has been cost. As

you can see (Fig. 23), chromium is a very attractive material for these

types of applications. Molybdenum has been changing drastically in cost

since 1975. Consequently, we have designed another group of steels with

low molybdenum to replace our high tonnage Equalloy applications. Since

1969 v;e have used these low nickel steels and then recently, the last

P7-12
- 12 -

couple years, we’ve switched over to a lower molybdenum type steel

referred to as Econalloy. As indicated, cost has been the motivating

factor ^nd, believe it or not, these small changes in alloying elements

amount to multi-million dollar a year savings in the tonnages we use.

But now I would like to show you som,e of the information that one has to

develop, engineering type information, to convince people that these

changes can be made. Referring to Fig. 24, the dotted line regions are

the LIO and L90 lines for failure due to profile degradation of gears and

the straight lines encapsulate the LIO and L90 for bending failure of

gears. This data represents fifteen years of testing of steels like

8620H. This represents how we think the standard performs in terms of

torque vs. cycles to failure. So on a new steel like the Equalloy or Ex

15, 16 steels we generate this type of data for comparison. We use a

rig like that shown in Fig. 25. It is called a four-square dynamometer.

The dotted points in Fig. 24 show then how the new steel performs in

com^parison to the standard. Another type of test we run is a chassis

dynamometer test. Shown in Fig. 26 is a chassis dynamometer. This test

simulates the shock loads the rear axle might experience if the vehicle

was being "frogged" from a mud hole after being stuck. The performance

bogey was 40,000 cycles in this case. The new steel EX l8 was equivalent

to if not better than the old steel (refer Fig. 27). Now statistically

this approach may be in trouble, but they’re equivalent as far as I can

see. Fig. 28 shows data from four-square type testing for the new lower

moly steels that v;e’re just now introducing. In this instance we tested

at tv;o different carbon levels. . ou can see the performance was very

P7-13
- 13 -

similar. In addition to evaluating manufacturing characteristics such as

machinability and distortion we do vehicle tests at proving grounds and

in customer service. This should give a little bit of an idea of what we

have to go through to convince .the "Pass the Aspirin" group that

substitution can be done. We believe, within limits, that the

performance will be equivalent if the carbon content, microstructure and

residual stress are very close to what was in the parts to be replaced.

I'd hasten to add that we're certainly aware that there are limitations.

We have been working with low alloys and the changes we're making are

relatively small. One of the limitations is that we don't know exactly

how large of changes can be made and at what alloy level the alloy per se

begins to affect performance. We know that it does take some insight

when using the system. Changes shouldn't be made by novices. The areas

where we need additional knowledge relate to influence of alloy on

fracture toughness and fatigue. We have made fracture toughness evaluations on the steels I just described. For instance, they all, the 8600

and the Equal loy steels have a Ky^ value .in the high carbon region of

about 15.5 Ksi square root inches. I'm not sure exactly what that means

for a gear, but they're about the same. There's room for research,

consequently, one should proceed with caution and be sympathetic with

the "Pass the Aspirin" guy and generate the data that's necessary to

provide the confidence that allows moving from one steel to another.

P7-14
- 14 -

Now I’ve discussed replacement steels, where we talked about having

exactly the same hardenability characteristics but with a different

chemistry. Now I'd like to touch briefly on the concept of alternate

steels. Altei'nate steels arc when a substitution is made with a

different hardenability steel but the steel still is adequate for the

application. Now what this requires is that some Application Tailoring

be done, ATing as we call it. I am going to use as an example hypoid

pinions of the type shown in Fig. 29. These are the sophisticated type

of rear axle drive gears used in trucks. One of the things that we have

to design against for these applications is the possibility of the

degradation shown in Fig. 30 happening. In heavy loaded off-highway

applications, these gears are stressed quite highly. The thing that

causes that type of degradation is the stress gradient along the tooth

profile not being supported by enough subsurface strength (refer Fig.

31). So what has to be accomplished is to establish what strength

gradient is required, convert that into hardness gradient required and

then select the hardenability steel and carburizing combination that

will provide that critical strength. We have gone into a considerable

amount of study on these gears, including knowing exactly what cooling

rates are going to be in production heat treat. This we do by running

"mock" carburized samples of different steels of known hardenability and

determine what we call Jominy Equivalent Conditions, JEC. After

extensive studies, we know that we have to have a minimum of strength

gradient as shown here. Fig. 32, for these components. The next step is

to take this information and turn that into metallurgical terminology

required by the computer (refer Fi . 33). Djg requirements, for

P7-15
- 15 -

instance, of the base material 3*9 to 6.8 inches. That will give the

hardness range required in the core area. The Djq of 6.8 inches minimum

will give the hardenability needed to provide the hardness

characteristics specified in the high carbon region of the case. This

information is placed into the CH part of the computer then the computer

program develops the least cost alloy composition (Fig. 34). What I would

like to point out (refer Fig. 35) then is that here are base H-bands for

two different steels, the new design steels which we call IHRM 20 which

will m.ake this part and give us the strength gradient required, and the

older steel 4820H, a high nickel steel, that had lower hardenability and

lower carbon, would also make the part, but was much more expensive.

These steels will both make adequately performing parts but they're

entirely 'different chemistries and different hardenabilities . The new

steel chemistry is shown here (refer Fig. 36). This is what the computer

provides. Note we're down from nickel to the .20 and .40 nickel

level. As a matter of fact, for your interest, noting the molybdenum

content here of .25 to .35, we are now designing low molybdenum steel

because of the increase in the price of molybdenum. It will be a CH

steel since it will have the exact same hardenability, it'll be a

replacement steel. We've used this steel for about five years now

successfully. Fig. 37 is a summary of what this did for the Company,

saved half a million dollars, cut down carburizing time and conserved

90,000 pounds of nickel per year.

That's the end of my story. In summ.ary, I'd like to say I've discussed a

technology that with some care in oplication can develop replacement

P7-16
- 16 -

steel ^ engineers and metallurgists work together they can develop

alternate steels. This technology is applicable to the chromium

situations. Chromium being a very attractive alloy in terms of

hardenability and cost will always be in these types of steel but the

restraint can be placed on the computer to keep it out and if one does

this it'll end up with steels that are more expensive but they would be

least cost without chromium. They would probably contain significant

amounts of molybdenum which is available on a domestic basis. I've

pointed out that the engineering commiUnity, the users of these steels, do

not favor making substitutions and that it's necessary to prove up the

steels in engineering and manufacturing type tests before they can be

accepted

.

I thank you for your attention.

P7-17
IDENTIFIED WORLD CHROMITE ORE RESOURCES*

(THOUSAND SHORT TONS) high Cr, HIGH Fe, & HIGH Al COUNTRY RESERVES OTHER

WESTERN HEMISPHERE UNITED STATES BRAZIL CANADA

CUBA GREENLAND

OTHER

50 6,800

280

6,100

5,750

2.900

1 ,200

1

1 ,000 200

HEMISPHERE TOTAL 7,130 27.150

WORLD TOTAL 1,866,000 2,965.000

' ADAPTED FROM US GEOLOGICAL SURVEY, PROF. PAPER 820, 1973

1

IDENTIFIED WORLD CHROMITE ORE RESOURCES

COUNTRY

HIGH Cr, HIGH Fe, t HIGH Al RESERVES OTHER EASTERN HEMISPHERE

FINLAND 11,000 5,600 GREECE 100 100 INDIA 7,800 6,700

IRAN 1 ,700 1,100 MADAGASCAR 5,600 5,600 PHILLIPPINES 5,280 2,760 SOUTHERN RHODESIA 616,000 616,000 SOUTH AFRICA 1,156,000 2,256,000 TURKEY 5,600 5,600 USSR 23,100 2A,200 OTHER 2,200 2,200 HEMISPHERE TOTAL 1 ,83li,380 2,925.860 WORLD TOTAL 1,866,000 2.965.000

' ADAPTED FROM US GEOLOGICAL SURVEY, PROF. PAPER 820,. 1973 2

1978 ALLOY STEEL PRODUCTION

AISI TONS ESTIMATED

SAE

Cr RANGE

FOR SAE TONS OF Cr REQUIRED

CLASS. MFG. GRADE GRADES %

Cr 1, '*72,913 5000-5100 . 70- 90 11,783.3 Cr-V 59,299 6100 .65- 95 1* 71*. 1

* Ni-Cr 137,2'i7 3100-9300 . 90-1 25 1,1*68.5 Ni-Cr-Mo-V 256,091 1*3BV00 .1*0- 60 1

, 280.5 Cr-Mo 2 , 687,512 1*100 . 80- 1 . 10 25,531 .1* Cr-Mo-V 1149,059 SPECIAL STEELS .60- 90 1

. 117.9

N

i -Cr-Mo 2,01*2,176 1*300-8600 .1*5- 65 11 , 231.9

S i-Mn 176,08'* 9200 .60- 80 1

, 232 . '6

TOTALS 6 , 980,381 = 66.8^

•' 51*, 120.5

4

C37-AA is: HRC 60 max at / c

J = 1/16 in.

C 1 ,5 mm) HRC 53 min at J - 1/16 in .

( 1 .5 mm)

DISTANCE FROM QUENCHED END, mm

10 15 20 25 30 35 40 45

12/4 6 8 10 12 14 16 1^20 22 24 26 28 30 32 DISTANCE FROM QUENCHED END, sixteenths In. / E7-15 means: ^ HRC 4 1 min at J = 7/16 in. (11 mm)

for 50% Martensite (0,37% C)

HRC 44 max atj = 15/16 m.

(24 mm) for 50% Martensite (0.44% C)

SUBSTITUTION:

(A) REPLACEMENT

(B) ALTERNATE

. 8640H hardenability band and classifying system (SponzilU,

Sperry, and Lyiell). 5 P7-18 6
HJSTOnr OF AUOV STEEL USAGE

7

Pass the Aspirin —

We’re Changing Materials

8

IH CHAT SYSTEM

Computer Harmonized Appncation Tailored

Quantitative System

For Design And Selection

Of Heat Treatable Alloy Steels

CHAT

APPLICATION' TAILORED

A FORMALIZED APPROACH USED TO DETERMINE

THE MINIMUM HARDENAB I LIT Y LEVEL REQUIRED

TO PRODUCE THE DESIRED HARDNESS AND STRENGTH PROPERTIES FOR A GIVEN APPLICATION

9

10

CHAT

COMPUTER HARMONIZED

COMPUTER SELECTION OF THE LEAST COST

STEEL ANALYSIS WHICH WILL MEET THE MINIMUM HARDENABILITY LEVEL CALLED FOR

BY APPLICATION TAILORING

1 1 P7-19
HARDENABILITY D , = MF^ x MF„ x MF^ . x MF„, x MF^ x MF,, 1C Mn Si Ni Cr Mo

Distance From Quenched End, 1/16 In.

13

0 0 2 0.4 0 6 0.8 Elemental Content, %

1 1 1 I I I j 0 2 4 6 8 10 12 14 16 18 20 22 24

Distance From Quenched End, 1/16 In.

18 17 P7-20
DISTANCE FROM QUENCHED END OF SPECIMEN IN MILLIMETERS

2 1 chemistry).

24
28

29 30

P7-22
< APPLICATION TAILORING. V ’ patannine OperaUonal Jee Range

^Severity and Variability of

4 Quench (J5-J9)

SalBase Caibon

r.28*.31%Cait>on

v'' '’Vax.CaraHaidness1lCS2 vaavi%« w a s^h sea aw9 a ^Dalculata HardanabHlty

- B (Cora 3^• 6^ Indies)

ic (Case 6S mches Min.) Select Caiburtzing Cycle

Check Siress/Strangth > CompaMbility

33 34

OPTUUZED COMPOSITION

^Reseaiehmaterial ^aHRMS20

ICafbon

^^Manganese .^.00-1.30%

iSidphur :£J02-JD4S

'^hoafilionis u:S35%Max. nsncon jao-i35%

^Chrandum :S0-1.95%

'flicfcei 30- jMK

SMybdenum .25-

P7-23

36
PRODUCTIONRELEASE

PROJECT SUMMARY

Aimual italMlal Savings of

HMOHonDolais

Savad 3860HomolCaitwtalng

FdmaooTkna

Consanad 3^,000 Cu. Ft Nanml Qas

Raducad Stratogte Alloy Usaga by

30,000 LBS/YR

Improved Steel MHl Yield and Billet Finishing Costs

37

P7-24
SUBSTITUTIONS FOR CHROniUM IN STEELS FOR COIL SPRINGS

S. T. Furr

Bethlehem Steel Corp.

P8

Laboratory Studies of LowChromium and Chromium-Free

Steels for Suspension

Coil Springs

SUBSTITUTIONS FOR CHROMIUM IN STEELS FOR COIL SPRINGS

S. T. Furr

Bethlehem Steel Corp. Bethlehem, PA

BECAUSE OF THE CHANGING political conditions

in the world, American industry must do

its homework well in advance to prepare

for the very real possibility that some

of the key alloying elements will be even

harder to obtain — and more costly —

then is now the case. Chromium is a good

example of an element with potential

supply problems, since 90% of the chromium

used in this country is imported from

South Africa, Turkey and the USSR. If the

supply of this important element becomes

tight, one of the tasks facing the steel

and automotive industries in this country

will be to move to spring steels in which

chromium can be replaced or reduced in content.

The need to develop a data bank of current information as well as new knowledge about alloy substitution in

heat-treatable steels was stressed at

a recent congressional hearing by Allen G. Gray, Technical Director of the American

Society of Metals (1)*. In addition to

permitting us to step in with substitutes

in the event of a tightened supply, such

knowledge could provide a spin-off benefit

in identifying more-efficient alloys for

noncrisis use.

STEELS FOR HOT-FORMED COIL SPRINGS

Over the years a number of alloys

have been used for hot-formed coil springs.

It will be useful to review these grades

since some may be potential substitutes

for the chromium-containing grades and

others may offer guidance in the design of new alloys. The chemical ranges of the

grades of most Interest are listed in

Table 1. Chromium-containing AISI

5160 (0.70-0.90% Cr) and modified 5160

(0.45-0.65% Cr) are at the top of the list

in terms of wide use, followed by three more alloy steels and two cart>on steels.

* Numbers in parentheses designate

references at end of paper.

ABSTRACT

AISI 5160 and modified 5160, the most

commonly used steels for automotive

suspension coil springs, contain chromium.

In the event this metal becomes hard to get, our industries should have the

information basis for introducing spring

steels with less or no chromium. This

paper reviews existing spring-steel grades

and summarizes a laboratory study aimed at finding new alloys with an optimum combination of properties. New alloy approaches

that are potentially competitive with 5160

and modified 5160 during normal times are Identified as well as substitutes that

could be used during an emergency.

0148 7191 /80/0225-0479$02.50

Copyright © 1980 Society of Automotive Engineers, Inc. P8-1
Except for AlSl 1095 and AlSl 6150,

the maximum carbon for these spring steels

is about 0.65%, which is apparently

considered the safe upper limit for spring

steels with good hardenabillty. Higher

Carbon will increase the danger of quench

cracking for those bars small enough

to through-harden.

The third grade listed in Table 1 is

AlSl 9260, a chromium-free grade containing

1,8-2. 2% silicon, that was commonly used

for automotive coil springs for many

years. In fact, both 9260 and a chromiumvanadium grade, 6150, were recommended in

the 1939 edition of the Metals Handbook

(2) for use in highly stressed applications.

While steels with a large silicon content,

as in the case of AISI 9260, have a reputation for excellent load-loss resistance, they may also have relatively poor

surface quality, i.e. , excessive seams,

pits and decarburization (3). To be able

to provide satisfactory fatigue life,

springs manufactured from as-rolled

bars must have good surface quality.

Modified 5160 is a successful compromise

combining the good load-loss resistance of the silicon steel (9260) with the good

surface quality of the chromium steel

(5160). The successful development of modified 5160 shows that a silicon addition of about 1% provides good load-loss

resistance without impairing surface

quality.

Another chromium-free steel was AlSl

4068 (4). As compared with 5160 (Table

1), the 4068 grade contains 0.20-0.30%

molybdenum rather than chromium.

Although AISI 1095, a carbon steel,

is not used for automotive suspension

springs, it is the most popular grade for

other hot-formed coil-spring applications

(5). One reason that alloy rather than

carbon steels are recommended for automotive applications is that the alloy steels

have better hardenabillty. Because of low

hardenabillty, quenched and tempered 1095

bars of various sizes show a wide range of yield strengths even though the surface

hardness is the same. Alloy steel bars

show less scatter since they through-harden.

For these reasons, 1095 is satisfactory

for coil springs manufactured from smalldiameter bars where hardenabillty is

adequate, or for springs manufactured from

large-diameter bars operating at low

stress where some hardenabillty can be

safely sacrificed (6). An example of a successful attempt to increase the hardenabillty of carbon

steel enough for automotive applications

is AISI 15B62 (7), Table 1. Increased

hardenabillty is obtained from a relatively

high manganese content and a boron

addition. AISI 15B62 was successfully

used during this decade as a substitute

for 5160, and its use was discontinued

only because of economic reasons. However, if other alloys for spring

steels become less available, AISI

15B62 could again be considered. The

technical success of 15B62 demonstrates

that the manganese level of spring steels

can be at least as high as 1.4%.

The logical next step after this brief

review of various spring steels that are

or have been on the market is to consider

how the requirements for spring steels

may change in the future. The most

important trend is the decreasing size of

automobiles. Spring steels for smaller

cars may need less hardenabillty and more load-loss resistance. The reason for less

hardenabillty is that the bar diameter

will continue to decrease as the size of

coil springs is scaled down in proportion

to the decrease in automobile size.

The need for spring steels with better

load-loss resistance could occur as a result of an effort to improve small-car

ride characteristics, which are inherently

rougher than those of large cars. There

are several ways to design a spring to soften the ride. One approach that may be

used in the future is to design variable-rate

springs (8). The, present approach is to decrease the spring rate in constant-rate

springs. However, the changes in coilspring geometry, i.e., larger coil diameter

and smaller bar diameter, needed to decrease the spring rate also increase the

operating stress and consequenaly the

danger of load loss. The problem ^of increased load loss due to decreased

spring rate was dealt with by adopting the

modified 5160 grade in 1977 as a replacement

for 5160 in springs with low spring rates. A sagging spring can pose a significant

problem by causing the bumper height to fall below the standard height required in

federal specifications (9).

TEST PROGRAM

Together with certain directions of interest expressed within the automotive

industry, this survey of spring steels

provided guidance in formulating the

objective and design of our test program.

The test program consisted of comparing

5160 and modified 5160 with other standard

P8-2
grades as well as experimental alloys in order to determine the best alloy-development

approaches for the design of low-chromium

or chromium-free spring steels.

The most important characteristics

required for spring steels are good

surface quality for satisfactory fatigue

life, good load-loss resistance to avoid

sagging, and sufficient hardennbil ity to through-harden the bar and thereby produce

a consistently high yield strength. The

alloys evaluated and method of analysis

are described below.

MATERIALS - The compositions of the

bars tested in this program are listed in

Table 2. These materials come from a variety of sources. Some were produced

from commercial heats and others from

laboratory heats. Some were tested during

the past year and others ten years ago. For example, the Code at the top of the

list, T54, represents bars rolled in 1969

from a 150-lb laboratory induction

furnace heat of 15B62. The third alloy.

Code C, represents bars rolled from a commercial open hearth heat of 1566 melted

in 1977. Some of the experimental heats

were fairly large. For example, the FI and F2 bars were rolled on a commercial

bar mill from experimental ten-ton electric

furnace heats in 1971. In general, the

bars studied in this program were less

than 3/4 inch (19 mm) in diameter.

However, experir.iental bars hot-rolled from

ox|5C r Imont al heats 01, 02, 05 and 06 were about 1-1/8 inch (29 mm) in diameter.

The first five codes listed in Table 2 are chromium-free carbon steels.

AISl 15B62 (Codes T54 and 33) is the only

standard spring steel in the group. As compared with 15B62, 1566 (Code C) contains

less manganese and silicon and no boron.

Codes 01 and 02, which are experimental

alloys containing copper, have more manganese than 15B62 but about the same amount of silicon. The primary reason for

adding copper to a carbon steel is to obtain Improved resistance to atmospheric

corrosion. Code 02 actually contains

slightly more manganese and silicon than

are allowed for a carbon grade. The

maximum limits for a carbon steel are

1.65% Mn and 0.60% Si.

The rest of the materials listed in

Table 2 are alloy steels. There is a good representation of standard grades —

four samples of 5160, one of 6150, three

of 9260, and three of modified 5160.

Codes FI and F2 are experimental 1% Si

steels similar in composition to modified

5160, except that the chromium level is

reduced from about 0.55 to 0.35%. Code F2 also contains a higher carbon content

(0.80%) than is typical for an alloy

steel. Codes F and S are experimental

alloys containing small chromium and

TABLt 2 - Chemical Composition of Steels Tested in Tliis ProKt‘^r^

C(ide Grade C Mn p s Si Cr AL Other Elements

T54 13B(S2 0. 60 1. 20 0. 013 0. 013 0. 33 0. 03 0. 003 0. 003

33 15B62 0. 60 1

.

23 0. 010 0. 021 n. 30 0. 06 0. 041 0. 001 ti

C 15hf) U. 60 1

.

0 ( 0. 012 0. 026 0. 20 0. 09 0. 02 7 -

01 Exp

.

0. 37 1. 30 0. 017 0. 01 i 0. 48 <0. 01 n. 013 0. 25 Cu

02 Rxp

.

0. 66 1

.

71 0. 016 0. 01 3 0. 68 <0. 01 0. 022 0. 3 3 Cu. 0.016 N

33 5160 0. 61 0. 71 0. 013 0. 016 (1. 18 0. 77 0. 0 3 -

3 5160 0. 66 0. 91 0. 021 0. 02 7 0. 29 0. 84 0. 0 33 U 5160 0. 60 0. 92 0. 008 0. 019 0. 36 0. 87 0. 04 2 - A1 5160 0. 60 0. 79 0. 008 0. 017 0. 2 3 0. 79 f). 0 36 0. 05 Mo

IH 6150 0. 31 0. 71 0. 00 7 0. 023 (}. 27 0. 93 0. 032 0. 24 V

3i 4260 0. 62 0. 83 0. 003 0. 006 2. 02 0. 08 0. 08 - 3H 4266 0. 68 0. 84 0. 006 0. 003 2,All 0. 1- 0. (I 1 0., 04 Mo. 0.,13 .Si 9260 0. 63 1

.

00 0..01 3 0. 021 1 . 87 ri. 08 0. (137 -

A M‘-d. 5160 0. 61 0. 92 0. 016 0. 013 0. 93 0., 36 f). 040 0. 0 3 Mo

B Mod. 5160 0. 62 0., 87 0. 013 0. 013 0. 97 n. 34 (1, (111 -

H M.id. 5160 0. 38 0., 90 0,. OKI 0. 020 (1.. 93 0.. 3.. 0. (1)3 * H Kxp

.

0. 61 0 .91 0.. 009 0,.017 n. 97 0,. 13 0., 080 - F2 Kxp

.

0., 80 0..91 0..010 0,.013 0., 90 0,. 36 0..036 -

F Ixp 0,. 69 1 . 1 3 0,.012 n,.01 3 0,. 62 0 . 12 0.. 016 0. 1 ] Mo

S r.xp

.

0 .61 fl . 88 0,. 012 (}..012 0,.99 0,.09 0..022 0.,05 Mo

03 Exp

.

0 .67 1 0 . 020 ti,. 019 n . 9 7 <0 Oi 0.. 036 _

Of) Exp

.

0 . 6/t 1 3 3 0 .016 n . 016 1 . 1 1 <0 . 02 0.. 040 -

93 Kxp

.

0 . 89 1 . 18 0 .017 (1 . 009 1 .07 0 . 04 n . 02 3 -

P8-3
molybdenum additions. In addition to these elements, Code F contains about

0.60% Si and Code S contains about 1% Si. Codes 05, 06 and 95 are chromium-free

alloys also containing about 1% Si. Codes

05 and 06 are 0.65% C alloys with 1.44 and

1.55% Mn, respectively. Code 95 is a

0.

89% C alloy with 1.18% Mn. PROCEDURE - The steels were tested in

the laboratory for hardenability and

load-loss resistance but not for fatigue

life, because results from dynamic tests

of laboratory specimens do not necessarily

correlate with the results from the

dynamic tests of springs. The fatigue

life of springs is dependent on the

surface quality of the spring round, since

the maximum stress is at the surface.

Defects and decarburization on the surface

of the spring can be traced to steel mill,

spring-making processes and testing

environment, but a laboratory study cannot duplicate these mill, manufacturing and

service conditions. However, since our survey of spring steels had shown that

poor surface quality in production is, in most cases, associated with large amounts

of silicon, we were able to minimize the

potential risk of surface problems by

limiting the silicon content of the

experimental steels to about 1%. In the final analysis, the practicality of experimental grades that look

promising on the basis of laboratory

hardenability and load-loss tests must be

proven in actual tests of coil springs and

in-service trials.

Hardenability - Hardenability was determined with the Jominy test according

to the procedure outlined in SAE J406b.

The Jominy specimens were quenched from

1600 F (870 C). Test specimen size

depended on bar diameter. If bar diameter

was less than one inch (25.4 mm), a one-half size specimen was used. If bar

diameter was greater than one inch (25.4

mm), both one-half-size and standard

specimens were machined for testing. When

available, mill Jominy data from cast specimens were obtained for comparison

with laboratory results. The hardenability

parameter (J 55) used to compare the

various materials is the distance from the

quenched end of the Jominy to the position

on the bar where the hardness is 55 Rockwell C, i.e. , the as-quenched hardness

required at the center of a spring round

( 10 ). Load-Loss Resistance - The relative

potential load-loss resistance was determined via the Bauschinger torsion test

(11). The test specimen, shown in Figure

1, was prepared as follows. Specimens

were rough-machined 0.050 inch oversize

out of as-rolled bars. The oversize

specimens were austenitized at 1600 F (870

C) , then oil-quenched and tempered to 50

Rockwell C, the hardness specified for

automotive suspension springs. After oil

quenching and tempering, the specimens

were ground to size.

Torsion tests were peformed on a standard torsion testing machine with a 10,000 in-lb (1130 N-m) capacity. A torsion test is performed by simply

twisting the specimen at a relatively slow

speed. In this program the rate of twist was 15 degrees per minute. For some of

the specimens the shear stress at a three-degree offset was determined from

the torque-twist diagrams recorded during

testing (12). For a torsion specimen of

the dimensions given in Figure 1, a three-degree-offset yield strength in

shear is approximately equivalent to the

0.2% offset yield strength in tension.

The Bauschinger torsion test measures

the "Bauschinger Effect," i.e., the

weakening of metal when it is strained in

the reverse direction after a plastic

strain in a forward direction. The

conventional torsion test reveals the

torque-twist diagram, which is a measure

of the stress-strain curve of the material

in one direction of twist only. For this

study we measured the Bauschinger effect

using the "actual Bauschinger twist"

(ABT), which measures the difference in

strain between the forward-twist and

reverse-twist curves at 75% of the maximum

torque (Figure 2).

Relating Load-Loss Resistance to the

Bauschinger Torsion Test - Since springs

are energy-storage devices, spring steels

must have considerable resilience, i.e., a large energy-storage capacity. Resilience

is the ability of a material to recover

its original size and shape after deformation caused by stress. The most logical

parameter for measuring the resilience of

spring steels is the modulus of hyperelastic resilience, because this parameter measures the energy-storage capacity after

a preset. Since the final step in spring

making is cold setting, the focus is on resilience after plastic strain.

The Bauschinger torsion test serves

to predict the potential load-loss resistance of spring steels, since a larger

Bauschinger effect means increased hyperelastic resilience, hence improved

resistance to load loss. As shown in

Figure 3 from our earlier work (11), the

better spring steel, 9260, has a larger

Bauschinger effect than 5160.

Chishiraa (13) recently verified the

utility of the Bauschinger torsion test. He compared a high-silicon steel, SAE 9254

P8-4
SPECIMEN FOR THE BAUSCHINGER TORSION TEST ACTUAL BAUSCHINGER TWIST

(0635 cm)

Fig, 1 - Torsion specimen used in the

experimental program

Fig. 2 - Torque-twist diagram showing

how the actual Bauschinger twist is defined

(14), and Japanese SUP 9, which is similar

to 5160, in the torsion test as well as in

static and dynamic tests of coil springs.

The 9254, with its larger Bauschinger

effect, did Indeed show much better

load-loss resistance.

RESULTS AND DISCUSSION

THE EFFECT OF COMPOSITION ON LOAD-LOSS

RESISTANCE - In Figures 4 to 8 the ABT

curves for 5160 and 9260 from Figure 3 are compared with ABT data for modified 5160

and other standard and experimental

steels. For all of these steels the

Bauschinger effect Increases with increasing

plastic strain.

Fig. 3 - Comparison of the Bauschinger

effect of AISI 5160 and 9260 spring

steels

As shown in Figure 4, the ABT data for

modified 5160 fall between the ABT curves

for 5160 and 9260. This observation

correlates with service experience, since

modified 5160 is used in place of 5160 in

highly stressed springs susceptible to load loss. The ABT for these standard

spring steels increases with increasing

silicon content — 5160 at 0.18-0.36% Si,

modified 5160 at 0.93-0.97% Si, and 9260

at 1.87-2.07% Si. The inherent reason for

this relationship will be discussed

below.

The ABT values for the oS:her steels

are grouped according to potential loadloss resistance. For example, alloys

compared with 5160 and 9260 in Figure 5 have about the same potential load-loss

resistance as 5160. Alloys with larger

ABTs and even better potential load-loss

resistance are compared in Figures 6 to 8. The ABT data plotted in Figure 5 were determined for 1566 (Code C), 15B62 (Code

T54), 6150 (Code 18), and an experimental

carbon steel (Code 01). The silicon

range of these alloys is 0.18-0.53%, and

the carbon range is 0.51-0.63%. Alloys

with about the same carbon content have

about the same Bauschinger effect as long

as the silicon level is fairly low.

Figure 6 gives the ABT for alloys FI,

F and 02, similar in carbon content P8-5
TWIST,

degrees

TWIST,

degrees

ACTUAL BAUSCHINGER TWIST

Fig. 4 - Comparison of the 5160 and

9260 actual-Bauschinger-twist (ABT) curves from Fig. 3 with ABT data for

modified 5160 steels

ACTUAL BAUSCHINGER TWIST

Fig. 5 - Comparison of the 5160 and

9260 ABT curves with ABT data for 1566

(Code C), 15B62 (Code T54), 6150 (Code

18), and an experimental carbon

steel (Code 01)

ACTUAL BAUSCHINGER TWIST

Fig. 6 - Comparison of the 5160 and

9260 ABT curves with ABT data for three

alloys with the same potential loadloss resistance as modified 5160

(see Fig. 4)

ACTUAL BAUSCHINGER TWIST

Fig. 7 - Comparison of the 5160 and

9260 ABT curves with ABT data for three

alloys with almost the same potential

load-loss resistance as 9260

P8-6
ACTUAL BAUSCHINGER TWIST

Fig. 8 - Comparison of the 5160 and

9260 ABT curves with ABT data for two alloys with a slightly better potential

load-loss resistance than 9260

(0.61-0.69%) to the steels in Figure 5 but

with more silicon. Comparing Figures 4 and

6, one sees that these alloys have the same Bauschinger effect as the modified

5160 alloys. The Code FI and modified

5160 alloys contain about 1% Si and

0.35-0.56% Cr. Codes F and 02 contain

0.62 and 0.68% Si, respectively, and no more than 0.12% Cr. Given the evidence

that silicon enhances the Bauschinger

effect, the fact that 0.62-0.68% Si steels

with little or no chromium have the same ABT as 1% Si steels with chromium indicates

that chromium decreases the Bauschinger

effect

.

The adverse effect of chromium on the potential load-loss resistance of 1% Si steels is confirmed by the data

illustrated in Figure 7. Alloys 05, 06

and S, which have the same amount of silicon as in modified 5160 but little

or no chromium, exhibit an ABT larger

than that for modified 5160 and almost

as large as that for 9260, the spring

steel with the best load-loss resistance.

According to the data in Figure 8 for codes F2 and 95, another element that

Increases the Bauschinger effect is carbon. Alloy F2 contains 0.80% C and

alloy 95 contains 0.89% C. The ABTs for

these experimental steels are actually

slightly larger than that for 9260 (15).

The excellent potential load-loss resistance obtained by increasing the carbon

content explains, at least in part, the

popularity of high-carbon steels such as 1095 for coil-spring applications.

To summarize to this point, the

Bauschinger effect increases with

increasing plastic strain. In quenched

and tenpered steels three elements influence

the Bauschinger effect. Carbon and

silicon increase the Bauschinger effect,

but chromium in the range of 0.35-0.56%

reduces this effect in 1% Si alloys.

Since a larger Bauschinger effect means better load-loss resistance, these results

provide the basis for the design of new

spring steels. THE INHERENT CAUSE FOR DIFFERENCES

IN THE LOAD-LOSS RESISTANCE OF SPRING

^ STEELS - A simple model (16, 17) can be

used to explain the effect of microstructure on the Bauschinger effect and

thus account for the positive Influence of

plastic strain and of carbon and silicon

on load-loss resistance. The microstructure

of quenched and tempered steels consists

of strong, noncoherent cementlte particles

embedded in a ductile ferrite matrix. If the steel is subjected to plastic strain,

dislocations accumulate at the matrixcarbide interface. When the stress in the

forward direction causing the strain is

released, the dislocations produce microresidual back stresses in the matrix

balanced by larger forward stresses in the

carbides. The existence, after plastic

strain, of microresidual back stresses in

the ferrite matrix of a quenched and

tempered steel was demonstrated by Wilson

using X-ray diffraction (18).

Such interphase stresses are estab- » lished in a spring by presetting, the

final step in spring making. IThen the

spring is deflected in service (the same direction as the preset), th^ interphase

residual stresses reduce the actual

magnitude of the applied stress because

the applied and residual stresses are in

opposing directions. Consequently, if a coil spring is protected by large interphase microresidual stresses, it will show

relatively low load loss in service.

The relative intensity of the interphase residual stresses can be measured

with the Bauschinger torsion test. The

twist in the forward direction, equivalent

to the preset of a spring, produces the

back stresses. The twist in the reverse direction reveals the stresses, since

microresidual back stresses add to the

applied stress to cause yielding at a lower apparent stress than in the forward

direction. Thus, with a larger Bauschinger

P8-7
effect being a measure of a material's

ability to develop large mic rores idual

interphase stresses necessary for good

load-loss resistance, the Bauschinger

torsion test can be used to predict the

potential load-loss resistance of spring

steels.

The magnitude of the back stresses

depends on the number of dislocations

accumulated at the carbides and therefore

increases with increasing plastic strain

and a larger volume fraction of carbides

(16). This model explains our results

showing that the Bauschinger effect

increases with increasing plastic strain

and carbon content. The effect of silicon on back

stresses is explained by its effect on carbide size. This alloying element

slows down the growth of carbides

during tempering by affecting the

transformation processes (19,20).

According to Fisher's analysis (21),

a combination of decreasing carbide

size and a larger carbide volume

increases the magnitude of the

Interphase stresses.

We verified that the carbides in a 1% Si steel are finer than the carbides

in 5160 by inspecting extraction replicas

from samples quenched and tempered to about 50 Rockwell C. As can be seen from

the electron micrographs in Figure 9,

clouds of small carbides were extracted

from the 1% Si steel and relatively large

rod-shaped carbides were extracted from

5160. The carbides from both steels were identified as cementite on the basis of

X-ray diffraction patterns obtained from

a Deb/e-Sherrer camera. Although both

films showed the same diffraction pattern,

the films were different in appearance.

Because the carbides are smaller, the

X-ray film for the 1% Si steel showed

fewer lines that tended to be broader

and more diffuse (22).

The simple model explained above

and supported with metallographlc evidence

is consistent with the empirical results

AISI 5160 CODE 06

(0 23% Si) . (

1 .1 I % Si )

Fig. 9 - Electron micrographs of carbides

extracted from 5160 and Code 06. * Carbides were extracted with celluloseacetate replicas from specimens etched

for 45 sec. in nital. Original

magnification was 30,OOOX

obtained from the Bauschinger torsion

test and with the correlations between

the laboratory test and service ratings

of spring steels.

HARDENABILITY - The hardenabllities

of the most commonly used spring steels,

5160 and modified 5160, are compared in

Table 3. The J 55 for these steels was determined from one-half size and standard

(cast) specimens. The J 55 for one-halfsize specimens tends to be larger than

that for standard specimens, because the

cooling rate at the same position on the

two bars is different (23). However, a comparison of data from either specimen

shows that the hardenability of these

standard grades is about the same. The hardenability values bf ten alloys

with less chromium than 5160 or modified

5160 are shown in Table 4. The alloys are divided into three groups on the basis of

TABLE 3 - Hardenability of AISI 5160 and Mod. 5160

155, 1/16 in (mm)

Code Grade c Mn Si Cr Other Elements

1/2 Size

Spec imen Standard

Spec imen

5 5160 0.66 0.91 0.29 0.84 13-1/2 (21) _ D 5160 0.60 0.92 0.36 0.87 - 11-1/2 (18) 8-1/2 (14) * A1 5160 0.60 0.79 0.23 0. 79 0.05 Mo 13 (21) “

A Mod. 5160 0.61 0.92 0.93 0. 56 0.03 Mo 22 (35) >8 (>13)

B Mod. 3160 0.62 0.87 0.97 0. 54 - 14 (22) 7-1/2 (12) *

H Mod. 5160 0.58 0.90 0.95 0.54 - 13 (21) 9 (14) it

* Cast P8-8
lABLt U H.i rdin.ih 1 1 i I y of Low-Clircnl um ind Ctiromiiim-Km’ Alloys

Poti-nt i.il Lt'.ui-Loss Code Cr.ide kfsisi.iru'i* c Hn

C 1 5f>6 Aht'u t O.bO 1.03 0.20

01 Fxp

.

S.inu* .'IP 0. 37 1.30 0. 48

33 l5Rf)2 SlhO 0. hO 1.23 0.50

FI Kxp

.

Ahi'ut 0. hi 0.91 0.97

F Fxp

.

S.ime ii.s 0.h9 1.13 0.h2

02 Exp. M.»d. SI 60 O.hh 1.71 0.h8

S Exp

.

Bc-t tt*r O.hl 0.88 0.99

03 Exp. TliiJn 0.h7 1 . UU 0.97

Oh Exp

.

Mod. SI 60 0. hA 1. 33 1.11 F2 £xp

.

0.80 0.91 0.90

* Cast

TABLE 5 - Yield Strength

Tempering

Temperature

Code Grade Base F (C)

C 1566 Carbon 680 (360)

01 Exp

.

Carbon 680 (360)

02 Exp. Carbon 74S (396)

A1 5160 Alloy 765 (407)

H Mod. 5160 Alloy 780 (416)

F Exp. Alloy 765 (407)

05 Exp. A1 loy 780 (416)

06 Exp

.

A1 loy 790 (421)

potential load-loss resistance, and within

each group are listed in the order of increasing hardenability. For example,

the first group consists of three alloys,

1566 (Code C), an experimental alloy (Code

01), and 15B62 (Code 33), that have about

the same load-loss resistance as 5160. In the case of 15B62, Code 33 was used to determine hardenability and Code T54 to determine load-loss resistance. The

hardenability of 15B62 is better than that

of Code 01, which, in turn, has better

hardenability than 1566. A comparison of

Tables 3 and 4 shows that 15B62 has about

the same hardenability as 5160 and modified

5160.

The next group of alloys. Codes FI, F and 02, has about the same potential

load-loss resistance as modified 5160.

The third group. Codes S, 05, 06 and F2,

has even better potential load-loss

resistance. The hardenability of these

steels is about the same or slightly less

than those of the currently used spring

steels.

J 5 S , 1/1 6 In (tnm) Other 1/2 Size .St.indard Cr K 1 fmont s Specimen Spec linen

0.09 3-1/2 (6) 3 (3)* <0.01 0.23 Cu 5-1/2 (9) 5 (8) O.Oh 0.001 B - 8 (13)

0. 35 - 11-1/2 08)

0.12 0.11 Ho 11-1/2 (18) - <0.01 0.33 Cu, 0.016 N 14 (22) 11 (18)

0.09 0.05 Mo 7 (11) <0.01 - 7-1/2 (12) 6-1/2 (10) <0.02 - 10-1/2 (17) 8-1/2 (14) 0.36 - 11-1/2 (18) -

of Carbon and Alloy Steels

Ha rdness

Rockwell C 3-deg Offset Yield Strength

psi (MPa)

49.7 127,700 (881)

50.0 131,200 (905)

49.7 128,400 (B85)

49.8 128,500 (886)

50.4 130,800 (902)

49.7 129,200 (891)

49.8 127,300 (878)

49.4 128,300 (885)

Because the trend in usage is toward

spring rounds with smaller diameters,

alloys with less hardenability than 5160

and modified 5160 are potential substitutes

for these chromium-containing grades.

YIELD STRENGTH - One reason for heattreating coil springs to high hardness is to obtain high yield strength. As can be seen in Table 5, a number of carbon

and alloy steels do not differ significantly from one another in yield strength,

when heat treated to about the same surface hardness. Presumably, all

of these steels had enough hardenability

to through-harden our test specimens.

Yield strength is an important

parameter for spring applications, because

it is a measure of the modulus of elastic

resilience, which is the Internal energy

to stress the material to the proportional

limit. In contrast to the modulus of

hyperelastic resilience, which correlates

with the Bauschinger effect, the modulus

of elastic resilience measures the energyP8-9
storage capacity of a material before

plastic strain.

Although high yield strength means

high resilience, experience has shown that

yield strength cannot by Itself rate various spring steels for load-loss

resistance. For example, 5160 (Code Al)

and modified 5160 (Code H) have the same yield strength but different service

ratings (Table 5). Given the problem that

ratings cannot be based directly on yield

strength data, it is useful it have a method such as the Bauschlnger torsion

test that can be translated into load-loss

res 1st ance.

LOW-CHROMIUM AND CHROMIUM-FREE SPRING STEELS

The selection of replacements for the

standard chromium-containing grades

depends on the requirements of the spring

application and on the availability of

chromium. For example, if the supply is

short and action must be taken quickly,

the most logical substitutes are 15B62 for

5160 and a grade such as FI for modified

5160 or 5160. The advantage of 15B62 is that it is already a proven grade for

automotive suspension coil springs. Our laboratory studies correlate with this

service experience by showing that 15B62

(Codes T54 and 33) has the same potential

load-loss resistance as 5160 (Figure 5)

and about the same hardenabil Ity (Tables 3 and A).

Code FI is an attractive replacement,

because the similarity of its composition

to modified 5160 makes it a low-risk

substitute. The laboratory work supports

this selection by showing that FI, with

about 0.2% less Cr, has the same potential

load-loss resistance (Figures A and 6) and

only slightly less hardenability (Tables 3 and A) than modified 5160.

Chemical ranges of 15B62 as a substitute for 5160, and of a grade based on FI as a substitute for modified 5160 or 5160 are as follows:

Tabulation A. Ready Substitutes

C

15B62 0.55

0.66

FI-Type 0.56

0.6A

tta ^

1.10 O.AO

l.AO 0.60

0.75 0.70

1.00 1.10

Cr B min

0.0005

0.25

O.AO

If work on replacement alloys begins

while the supply of chromium is still

adequate, there will be time to perform

actual trials of experimental alloys to find the best new steels. The first step

is to decide how much hardenability and

load-loss resistance are needed. The loadloss requirement will depend on the spring

design. The amount of hardenability will

depend on the size of the round and severity

of the quench (10).

Current service experience is a good

guide to the amount of load-loss resistance

needed. For example, if the load-loss

resistance of 5160 is good enough for

the application, the best choice for a replacement could be a carbon steel

similar to Codes 01 and 02, since these

experimental grades have about the same or slightly better potential load-loss

resistance than 5160 (Figures 5 and

6). If more load-loss resistance is needed, the best choice could be an alloy steel with about 1% Si, similar

to Codes 05 and 06. These alloys have a better potential load-loss resistance than

modified 5160 (Figures A and 7). The main

source of hardenability in these steels is ^manganese with some help from silicon.

Assuming that the maximum possible hardenability is required for the application,

the chemical ranges would be as follows:

Tabulation B. Projected Substitutes

C Mn Si Cu

Carbon

Alloy

0.57

0.67

0.56

0.6A

1.35

1.65

1.30

1.60

O.AO

0.60

0. 70 1.10

0.20

O.AO

If less hardenability is required, less

manganese would be needed.

Although the steels in Tabulation A

could quickly move into use if there is a chromium availability problem, those in

Tabulation B would, of course, have to be

proven in actual tests of coil springs.

ACKNOIVLEDGEMENTS

The author would like to thank his

colleagues at Bethlehem Steel for many

helpful discussions. He is particularly

indebted to J. J. McCarthy for the design

of the new carbon grade. B. S. Mikofsky's

help in manuscript preparation is gratefully

acknowledged.

REFERENCES

1. A. G. Gray, "Materials Policy,

Research and Development Act of 1979

(H.R. 27A3) Including Comments on Substitution Technology for Critical

Materials," Testimony at Hearings of

Subcommittee on Natural Resources and

Environment in Joint Action with the

P8-10
Subcommittee on Science, Research and

Technology of the Committee on Science

and Technology of the U.S. House of

Representatives, June 28, 1979.

2. Metals Handbook, 1939 Edition,

ASM, "Composition and Heat Treatment

of Large Helical Springs," pp. 985-986.

3. E. T. Bittner, "Alloy Spring

Steels," Trans. ASM, Vol. 40, 1948,

pp. 263-280.

4. Metals Handbook, Edited by

T. Lyman, 1948 Edition, ASM,

"The Composition and Heat Treatment

of Steel Springs," pp. 640-647.

5. Metals Handbook, Edited by

T. Lyman, Eighth Edition, ASM, Metals

Park, Ohio (1961), "Steel Springs,"

pp. 170-171.

6. Metals Handbook, Edited by

B. P. Bardes , Ninth Edition, ASM, Metals

Park, Ohio (1978), "Steel Springs,"

(J. H. Maker), pp. 297-303.

7. R. F. Kern, "Carbon Steel

Spring Elements," U.S. Patent 3,689,329

(Sept. 5, 1972). Assigned to Caterpillar Tractor Company.

8. A. R. Solomon, "Coil Spring

Design, An Analytical View," Paper 790410

presented at SAE Congress and Exposition,

Detroit, March 1979.

9. F. Borik, V. A. Biss, and

Y. E. Smith, "Sag Resistance of Si-Mo

and Si-Cr Spring Steels," Paper 790409

presented at SAE Congress and

Exposition, Detroit, March 1979.

10. R. F. Kern, "Selecting Steels

for Heat-Treated Parts: Part 1 - Direct

Hardenable Grades," Metal Progress,

Vol. 94, Nov. 1968, pp. 61-73.

11. S. T. Furr, "Development of a New Laboratory Test Method for Spring

Steels," Journal of Basic Engineering,

ASME Trans. Vol. 94, Series D, 1972,

pp. 223-227.

12. G. E. Dieter, "Mechanical

Metallurgy," McGraw-Hill, New York

(1961), pp. 273-281.

13. K. Chishima, M. Ohhara and

K. Uchibori, "On Anti-Creep Property

of High Stress Coil Springs Made of SAE 9254," Paper 10 presented at JSSR

(Society of Spring Research of Japan),

Osaka, Japan, Nov, 1978.

14. F. P. Zlmmerli, Discussion

of "An Investigation of Tempered

Chromium-Silicon Spring Steel," by

H. J. Elmendorf, Trans. ASM, Vol. 40,

1948, pp. 298-299.

15. S. T. Furr, "Helical Spring

Steel and Method," U.S. Patent 3,847,678

(Nov. 12, 1974). Assigned to Bethlehem

Steel Corporation.

16. R. M. Jamieson and J. E. Hood,

"Bauschinger Effect in High-Strength LowAlloy Steels," Journal of the Iron and

Steel Institute, Vol. 209, 1971,

pp. 46-48.

17. D. V. Wilson, "Reversible Work

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Acta Met allurgica , Vol. 13, 1965,

pp. 807-814.

18. D. V. Wilson and Y. A. Konnan,

"Work Hardening in a Steel Containing a Coarse Dispersion of Cementite Particles,"

Acta Metallurgica, Vol. 12, 1964,

pp. 617-628.

19. J. Gordine and 1. Codd, "The

Influence of Silicon up to 1.5 wt-% on

the Tempering Characteristics of a Spring

Steel," Journal of the Iron and Steel

Institute, Vol. 207, 1969, pp. 401-407.

20. R. M. Hobbs, G. W. Lorimer and

N. Ridley, "Effect of Silicon on the

Microstructure of Quenched and Tempered

Medium-Carbon Steels," Journal of the Iron and Steel Institute, Vol. 210, 1972,

pp. 757-764.

21. J. C. Fisher, E. W. Hart and

R. H. Pry, "The Hardening of Metal

Crystals by Precipitate Particles,"

Acta Metallurgica, Vol. 1, 1963,

pp. 336-339.

22. B. D. Cullity, "Elements of

X-Ray Diffraction," Addison-Wesley

,

Reading, Massachusetts (1956), pp. 93-99.

23. Isothermal Transformation Diagrams

Third Edition, United States Steel,

Pittsburgh, Pennsylvania (1963), pp. 172-173

P8-11
Ai\l UPDATE ON FE-MN-AL STEELS AS A REPLACEMENT

FOR NI-CR STAINLESS STEELS

Samir K. Banerj

i

Foote Mineral Co.

0

P9
AN UPDATE ON FE-I-IN-AL STEELS*

FASRAD 81-5

Samir K. Banerji

Foote Mineral Co.

Route 100

Exton, PA. 19341

Introduction

An idea was put forth by the author a fevj years ago to look into the

potentials of an Fe-Mn-Al alloy system as a replacement for the austenitic

Ni-Cr stainless steels in some non-critical applications (1,2). That generated

quite a bit of interest around the world and even sparked some early investigations in some laboratories. Preliminary results from some of these investigations and information from some additionul literature which has come to my

r attention since the publication of my earlier papers (1,2) are sximmarized here.

However, before proceeding to that a brief account of the criticality of the

chromium situation is presented to empliasize the relevance of pursuing this

idea of developing austenitic stainless steels without chromium and nickel.

Although the situation regarding Ni is not so critical, its replacement by Mn

is feasible and attractive from an economic and weight-reduction viewpoint.

Criticality of the Chromium Situation

Our vulnerability regarding the chromium situation is well known. A

comprehensive review on the "Contingency Plans for Chromium Utilization" has

been published by the National Materials Advisory Board (NMAB) of the National

Academy of Sciences (3). U.S. is totally dependent on imports for all its

chromium needs. About 95% of the world reserves of the chromium ore are

located in only two countries: South Africa and Rhodesia (see Table I). The

NMAB report cited above (3) states that in the past, U.S. has imported its

chromium from several widely separated geographic locations (see Table II)

but it is estimated that these secondary sources may soon be exhausted or *To be presented at Workshop on "Conservation and Substitution Technology

for Critical Materials", Nashville, Tennessee, June 15-17, 1981

P9-1
2

become insignificant. U.S. chromium deposits are small and virtually no

prospects exist for the discovery of any significant new U.S. deposits. There

is little evidence to suggest the existence of significant chromium deposits

outside of those geographic areas presently known. Therefore, it is feared

that at the current rate of consumption, within 25-75 years the world will

depend completely on South Africa and Rhodesia. This is an awesome fact and

yet we have done very little in terms of substitution, conservation, reclaimation and recycling of chromium.

It is obvious that a chromium embargo or a cartel will virtually strangle

the nation and pose a grave threat to the national security because "no substitutes exist or are likely to be developed for chromium in the high-strength

steels, high-temperature metals, and corrosion-resisting alloys that are essential in the manufacture of jet engines, petrochemical and power plant equipment,

and various other critical products (3)". This constitutes about 35% of the

total U.S. chromium usage which is essential and, therefore, the remaining 65%

of all the chromium used is actually non-essential and replaceable . Figure 1

shows that about 65-70% of the total U.S. chromium use is in metallurgical

applications and the remainder being in refractories and chemicals. Table 3

shows the details of Cr consumption under each major use classification and also

the possibilities of savings in each application. This Table also shows that

it is technically feasible to replace about 26% of the total U.S. chromium use

almost immediately and another A0% can be saved after 10 years of research and

development. A crash program could obviously accelerate this future potential

saving

.

As indicated in Table 3 the largest single use of chromium is in stainless

and heat-resisting steels, constituting greater than 50% of all chromium consumption in the U.S. Therefore, conservation and substitution of chromium in

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3

this category will provide the maximum savings in terms of tonnage. About

40% of chromium used in these steels is irreplaceable but about 60% of it can

be replaced. It is technically feasible to replace about half of that (30%)

chromium almost immediately. However, very little has actually been done in

this respect. It should be realized by all that in a large number of applications of stainless steels they are "over-designed" for the purpose. This is

so because we live in a wasteful society and have been spoiled by the ready

availability of chromium so far at a relatively low cost. This plus the outstanding versatility of austenitic stainless steels have made us choose this

path of least resistance. Chromium-free or low-chromium substitutes could

easily be used in decorative applications, automotive trims, flatware, kitchen

utensils, sporting goods, home applicances etc. where traditional high-chromium

stainless steels are used unnecessarily. A recent article by Watson (4) states

"we could dispense with stainless steel sinks in our homes, and use aluminum

in place of stainless steel for our pots and pans. Stainless steel used in

dairies, food preparation plants and restaurant kitchens as well as for the

transportation of food could be replaced, but not easily or cheaply". Such

ideas have to propagate widely and expeditiously. Certainly, no substitution

or conservation program is going to be easy or cheap but knowing the criticality

of the chromium situation every effort must be made to save any and every

ounce of chromium from non-essential applications to direct it toward those

essential uses where chromium is absolutely irreplaceable. Table 4 shows the

various end uses of stainless steels and the chromium savings possible in each

category. Note that chromium-containing stainless steels used in ordnance and

accessories, mining, petroleum, natural gas etc. are irreplaceable but in most

other categories large fractions of the chromium use can be replaced.

Some of the technologies available for immediate substitution and conP9-3
4

servation of chromium in stainless and heat-resisting steels are, for example,

clad and coated products, replacement of higher Cr steels with existing lower

Cr stainless steels in non-critical applications, and using alternate materials

such as aluminum, titanium, zinc, copper, brass, plastics etc. wherever

possible. As mentioned earlier such measures could immediately save about 30%

of the Cr used in stainless steels. The other 30% of the replaceable Cr in

stainless and heat-resisting steels will require some research and development

efforts as indicated in Table 3. These efforts will include developing chromium

free or low-chromium substitute steels. Again, it must be reiterated that such

substitute materials may be good only in non-severe or non-critical applications

It should also be realized that such substitutes may not be as universal and

versatile as the present day high chromium stainless steels and one may need

to develop specific alloys for specific purposes. The remainder of this paper

will deal with one such idea of a chromium-free alloy which appears to have

good potentials of replacing austenitic stainless steels in some non-critical

applications

.

The Fe-Mn-Al Alloys

In previous papers (1,2) it was shown that it is feasible to replace

chromium with aluminum, and nickel with manganese and carbon, and yet produce

an austenitic alloy, with better mechanical properties than the conventional

Ni~Cr austenitic stainless steels. Such alloys can have good oxidation and

sulfidation resistance and moderate corrosion resistance. Since the publication of those papers several other references were found in the literature on

this subject. These are briefly summarized below.

Summary of Additional Literature on the Fe-Mn-Al System

1) Fe-Mn-Al Precipitation - Hardening Austenitic Alloys by G. L. Kayak (ref. 5)

This and the subsequent several Russian papers deal with the "Fermanal”

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5

or 9G28Yu9 (0. 9%C, 28-29%Mn, 9%A1) type steels with or without Mo, W, and Cb(Nb)

which are about 12-15% lighter than common steels with a density of about

6. 7-6.8 g/cm^. They discuss the mechanical and magnetic properties, aging

and precipitation kinetics, oxidation and some corrosion resistance of these

Fe-Mn-Al steels. This particular paper by Kayak (5) mentions use of these steels

in aviation and automotive industries for critical and heavily loaded parts

e.g. springs, gears, landing gears etc. He found an optimum composition of

25-30%Mn, 8-10%Al, about 1%C, bal. Fe to get an austenitic structure and good

physicochemical properties. Figure 2 shows the mechanical properties of such

an alloy as a function of aging temperature. A typical age-hardening peak is

observed around 550°C after 16 hours of aging. Increasing the aluminum content

and decreasing the manganese and carbon contents results in ferrite phase in

the microstructure impairing the plasticity and magnetic properties of the

alloy as reflected by permeability studies shown in Table 5. The increased

permeability after aging in alloys containing only 23.7%Mn or greater than 10%

A1 was believed to be due to the precipitation of iron-aluminum carbides and a

breakdown of the austenitic structure. The alloys with 28-30% Mn and about 8%

Al were completely non-magnetic and austenitic even after aging for 20 hours

at 600°C. The oxidation resistance of the alloy was found to depend mainly on

the aluminum concentration and very little on the manganese content, which actually impaired it slightly. Alloys with more than 8%A1 showed a weight gain

of 6-12.5 g/m^ in air at 700°C and 500 hours. The lKhl7N2 (17%Cr, 2%Ni) stainless steel and the 20KhM (Cr-Mo) boiler steel showed weight gains of 6.2 g/m^

and 118 g/m^, respectively under same conditions. This illustrated the good

oxidation resistance of these Fe-Mn-Al alloys.

2) Phase Composition, Structure, and Properties of Low-Density Steel 9G28Yu9MVB

by M. F. Alekseenko et. al. (ref. 6)

P9-5
6

This paper discusses the development of a magnetic method for the quantitative phase analysis of "Fermanal” type alloys in the as-quenched condition

and after aging. It was determined that the age-hardening is caused by the

precipitation of a complex carbide of the (FejMn) 3AlCx type. This is consistent with the work of James (7) who used electrolytic extraction and x-ray

diffraction techniques for the identification of this phase. The amount of

this hardening phase was maximum at about 550°C aging temperature resulting

in a peak hardness (Fig. 3) and strength (5) and minimum ductility (Fig. 2,

ref. 5) and impact strength, Fig. 3. Above this aging temperature they believed

that the (Fe,Mn)3AlCx phase starts to dissolve causing the strength to drop and

ductility to increase. However, the toughness may or may not recover because

it depends not only on the amount of the hardening phase but also on the grain

size and the presence and size of precipitates at the grain boundaries. Based

on this study these authors recommended that for optimum mechanical properties

the A1 content of the alloy should be limited to a maximum of 9.2% with carbon

not exceeding 0.92% and a fine-grained fully austenitic structure should be

obtained after quenching.

The work of James (7) on a Fe-9Al-9.5Mn-0.96C-6.77Ni alloy also showed

similar precipitation of a mixed carbide phase (FeMn) 3AlC with a cubic

structure. This precipitation occurred in a cellular form at austenite grain

boundaries as well as twin boundaries. The peak hardening in this case occurred

at about 600°C aging temperature in both the as-rolled and solution treated

condition. Fig. 4a. Kinetically, a sharp increase in hardness occurred after

only 15 minutes of aging with the peak hardness achieved in 2 hours for the

as-rolled material compared to 4 hours in the solution-annealed condition.

Fig. 4b. This alloy showed a sevei^e embrittlement as a result of aging and

the gain in strength was offset by the loss in ductility and toughness.

P9-6
7

3) Phase Transformation Kinetics in Steel 9G28Yu9MVB by G. S. Krivonogov

et. al» (ref. 8)

This paper takes a deeper look into the kinetics of austenite decomposition

in these steels as a result of aging over a wide range of time and temperature.

Quantitative magnetic methods for phase analysis, electron microscopy and x-ray

crystallography are used to show that the decomposition of the as-quenched austenite occurs in two stages upon aging. The first step (at aging temperatures

below 700°C) involves the precipitation of a finely dispersed coherent K-phase at

450-550°C which has an ordered f.c.c. structure with Fe 3AlCo .65 composition and

is oriented in the <100> crystallographic direction. This is the dispersion hardening phase which can be used to achieve the required combination of mechanical

properties with the matrix still being austenite. With increasing aging temperature these K-phase particles grow, their number density drops, and their coherency with the matrix is lost. Above 650°C large K-phase precipitates are observed only at the grain boundaries. The second stage occurs around 650-750*^0

when the austenite which is now depleted in carbon and aluminum and rich in

manganese breaks down into ferrite and a brittle beta-manganese phase. The

K-phase is simultaneously enriched in Mn and C by redistribution of the alloying

elements and can contain up to 15-25%Mn after prolonged aging at 650-750°C. This

is the (FeMn)3AlCx phase mentioned earlier (6). At aging temperatures of 750-

850°C this precipitate forms in the first stage of decomposition. In the second

stage, the beta-manganese is not formed at these temperatures, instead more of

the (FeMn) 3AlC2j phase precipitates and austenite transforms to ferrite rich in

Mn and Al. Prior plastic deformation of the austenite substantially accelerates

the second stage. Plastic deformation after aging destroys the ordered structure of the carbide phase and this type of treatment can be used to produce

ultra-high-strength sheets with tensile strength of 170-180 Kg/mm^.

4) Structural and Phase Changes in Steel 9G28Yu9MVB During Aging by L. I. Lysak

P9-7
8

et. al» (ref.

9

)

This is practically the same paper as ref. (8). It again shows the agehardening peak around 550°C being related to the amount of finely dispersed

coherent K-phase. The loss in coherency, growth, and reduced number of these

precipitates explains the decrease in hardness and strength at higher aging

temperatures. Plastic deformation of the quenched austenite again accelerates

the kinetics of aging.

5) Nature of Age Toughening of Fe-Mn-Al-C Alloys by N. A. Storchak and

A. G. Drachlnskaya (ref. 10)

This is again a very similar paper to references (8) and (9) . The same observations were confirmed by x-ray diffraction methods using single crystals of

the 9G28Yu9MVB steel grown by the recrystallization annealing method at 1000-1100°C.

6) State of the Surface Layer and Corrosion Resistance of Steel 9G28Yu9MVE by

V. P. Batrakov et. al. (ref. 11)

First, the mechanical properties of the cold-rolled sheets, and quenched

(from 1050°C) and aged (at 500-600°C for 10-20 hrs.) bars of this steel are provided. At normal test temperatures the cold-rolled sheets (50% reduction) have

a tensile strength of 140-150 Kg /mm^, 0.2% yield strength of 130-140 Kg/mm^

with 10-12% elongation. The bars have a tensile strength of 110-120 Kg/mm^,

0.2% yield strength of 100-105 Kg/mm^, 35% elongation, 55% reduction in area

and impact toughness, A^=7-9 Kgm/cm^. The stress-rupture strength, CJ^qq-»' 70

Kgm/ cm^ is reported. The use of this steel at temperatures up to 500°C is

recommended.

The general corrosion behavior was measured in wet, tropical, and salt

mist atmospheres. Specimens were degreased, etched, and passivated before

placing them in the test chamber. The wet chamber was maintained at 20-25°C

and 95-98% relative humidity. The tropical atmosphere test comprised of holding

P9-8
9

for 8 hours at 50°C and 95-98% relative humidity, 12 hours at 20+5°C and 95-98%

humidity, and 4 hours at 2C^5°C and air purging. The salt mist chamber was

held at 20-25°C with 95-98% relative humidity and periodic spray of a 3% NaCl

solution. Maximum corrosion occurred in the salt mist environment. The practice

of iron shot blasting to remove the scale after hot rolling and quenching before

cold rolling displayed friable corrosion products at the points of impingement after only 7 days. In this condition the steel showed an unacceptably

high corrosion rate (5-lOmm/yr . )

.

Elimination of shot blasting, and deep chemical etching to a depth of

100-300y to remove the scale in hot rolled and quenched sheets provides satisfactory corrosion rates (0. 01-0. Imm/yr . ) after cold rolling and passivation.

Examination of the surface layer formed in this steel when it is exposed to

high temperatures revealed the reasons for this corrosion behavior. Electron

microprobe analysis revealed a ferritic surface layer depleted in Mn and A1

by factors of 4 and 2, respectively compared to the bulk composition when the

steel is exposed at 1100°C for 1 hour. The thickness of this outer layer is

about 20-25 microns under these conditions, Fig. 5. Metallographic and magnetometric studies confirmed this outer layer to be ferritic and also showed

that this ferrite phase actually extends deep into the metal along the

austenite grain boundaries. Figure 6 shows a plot of the depth of penetration

of ferrite at the grain boundaries from the metal surface as a function of

time at three different exposure temperatures, determined by metallographic

and saturation magnetization (47TIS ) methods. Note that the depth of the twophase layer actually extends as deep as 350-400 microns after 1 hour exposure

at 1100°C or about 500 microns at 1150°C. The total ferrite content of the

specimen does not exceed 6-8% as estimated from the saturation magnetization.

Therefore, satisfactory corrosion resistance is attainable only if the twoP9-9
10

phase layer is removed completely, otherwise pitting corrosion occurs where

the ferritic grain boundaries intersect the outer surface.

For best results the scale and the two-phase surface layer of the hot

rolled and quenched steel should be etched away in 20% HNO3 at room temperature

for 30-40 minutes or by 40% KOH + 2% potassium permanganate at 120°C for 1-2

hours followed by 5 minutes in 20% HNO 3 at room temperature. After cold rolling

the steel should be pickled and passivated in a solution of CrO^ (150 g/liter) +

H

3PO4 (100 ml/liter). For more severe environments further protection against

corrosion can be achieved by depositing paint and varnish coatings on specially

prepared surface which can be attained by galvanic or phosphate coating, or

metal spraying with aluminum. For zinc or cadmium galvanic coating the surface

can be prepared by etching in 20% HNO3 for 2-3 minutes followed by chemical

degreasing and pickling in 4-5% solution of orthophosphoric acid for 30-90

seconds. Such systems of paint and varnish coatings based on epoxy-polyamide or

silicone enamels, combined with phosphating or spraying with aluminum ensure

protection from corrosions at operating temperatures up to 500°C. Again, it is

emphasized that for good adhesion of paint and varnish coatings it is essential

to remove the two-phase surface layer completely.

One very interesting feature of this steel is that it was found to be not

susceptible to stress corrosion cracking. Pure bend specimens stressed to 80%

of the yield stress were tested in the salt mist and tropical chambers. Some

specimens were only degreased and others were degreased, etched, and passivated.

No effect of stress on the overall corrosion was observed. Despite the presence

of pitting, no stress corrosion cracking was detected after 90 days of testing

in these atmospheres.

7) Aluminum-Manganese-Iron Alloys by R. E. Cairns, Jr. and J. L. Ham, ref. (12)

This 1963 patent is a much detailed work along the lines reported earlier

P9-10
11

by the authors (13) and referenced in my previous papers (1,2). The principal

objective of this invention was to provide low density, oxidation resistant,

cold-ductile austenitic alloys from non-strategic materials having good mechanical properties at room and elevated (1200°F) temperatures. It was also the

object of this work to achieve good workability so that the alloys can be

fabricated by normal steel processing techniques. These properties were

satisfactorily achieved in the patented alloys with following composition:

20-A0%Mn, 7-16%Al, 0.15-1.1%C, 0-10%Ni, and bal. substantially iron. The

forgeability of these alloys are reported to be very good and a wide temperature

range of 2000-2300°F could be used without harmful effects. The machinability

of the alloys containing more than 15% austenite was comparable to normal

stainless steels. The composition of some specific alloys, the percentage

of austenite in those alloys, and their mechanical properties at room temperature and 1200 F after different heat treatments are shown in Table 6. The

improvement in room temperature ductility with increasing austenite content in

the alloy is clearly shown in Fig. 7.

The density of these alloys are approximately 15% lower than those of

common stainless steels and much lower than those of high temperature superalloys. The oxidation resistance of these alloys in general are quite good as

shown in Table 7 and compared with a stainless steel (310) , a cobalt-base

alloy (S-816) , and a nickel base alloy (Udimet 500) . Note that increasing Mn

content does impair the oxidation resistance of the alloy somewhat. In cold

rolling no difficulty was encountered in 80-100% austenitic alloys and normal

steel processing techniques could be used. 80% cold reduction was achieved

with 20% per pass. The alloy containing A0% austenite could be cold reduced

66% using the above schedule. The strain-aging characteristics of these alloys

were shown by annealing the cold rolled sheets (0.1") for 1 hour in the temperP9-11
12

ature range 700 -1000°C after different degrees of cold reduction. A minimum of about 20% cold reduction is necessary to observe the strain—agehardening phenomenon which occurs around 700°C annealing temperature, for

example, see Fig. 8 for alloy #213 containing 10.2%A1, 34.4%Mn and 0.76%C.

Therefore, the mechanical properties shown in Table 6 really reflect the solution-annealed properties and not the maximum attainable values.

8) Iron Aluminum Alloys by W. H. Richardson (ref. 14)

This 1965 patent describes a family of Fe-Mn-Al alloys very similar to

those patented by Cairns and Ham (12). The object here being the development

of alloys possessing good ductility, high strength, corrosion resistance and

low density, combined with high electrical resistivity and low magnetic permeability. The range of composition of these alloys are; 18-40%Mn, 4-14%Al,

0. 15-2%C, with at least one of the following additional elements 0.5-12%W,

0.1-5%Mo, 0.5-15%Ni, 0.5-5%Co, 0.1-2%Cb, 0.5-15. 9%Cr, up to 5%Ti, Cu, V, or

Zr, or up to 1% B or Ce, up to 1.5% impurities (P,S,Si), the balance being iron.

These alloys can be produced by the conventional electric furnace steelmaking

methods or by induction melting. Both continuous casting or conventional ingot

casting can be used. The ''Durville'' casting process gave excellent result.

The preferred soaking temperature of ingots or billets prior to hot working

is recommended to be 1100°-1200°C for one-half hour per inch of thickness or

cross sectional area. They can then be hot worked satisfactorily by any of

the conventional hot working processes such as rolling, forging or extruding.

Table 8 shows the various alloys used in this work. The mechanical properties

of hot-rolled (from 1100-1200°C) bars were measured. The ductility and impact

properties could be improved considerably by a solution-annealing treatment at

800°-1250°C after hot rolling. Alternatively, similar improvement in properties

could be achieved by finishing the hot working operation in the solution

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13

annealing temperature range and then quenching from that temperature, thus,

eliminating the need for subsequent solution anneal. Further hardening can

be produced by aging at 400-700°C and any combination of strength and toughness

properties can be achieved by adjusting the time and temperature of aging.

The solution annealed (at 800-1250°C) and quenched alloys are amenable to

cold working and tensile strengths higher than 200,000 psi are achievable after

40-50% cold reduction. The cold rolled materials also respond to strain-age

hardening in the temperature range of 400°-700°C. In this conditions, hardness

values as high as 600-700 V.P.N. may be achieved making the alloys suitable for

severe wear and abrasion resistant applications. Intermittent corrosion tests

2 in 5% NaCl solution showed a corrosion rate of 49.8 mg/cm /day in a 15 day test

for an Fe-8.6Al-19.7Mn-l.02C allqy. A 7 day test on an Fe-6.lAl-30.6Mn-0.56C

alloy showed a corrosion rate of 60.1 mg/cm'^/day. Addition of Cr and/or Ni

to such alloys reduced the corrosion rates by about an order of magnitude.

Alloys of this invention containing less than 7%A1 are not susceptible to embrittlement during prolonged aging at 500-600°C and they also possess very

good impact properties at sub-zero temperatures.

These alloys could be modified by adding some Ni and dropping the Mn but

still maintaining the predominantly austenitic structure. Such modified alloys

could contain 4-14%Al, 8-18%Mn, 3-15%Ni, 0.6-1.2%C, bal. Fe. The mechanical

properties and heat treatment response of these Ni-modified, low-Mn steels are

quite similar to the Ni-free, high-Mn steels described above. All these steels

are predominantly austenitic but may contain a small amount of ferrite depending

on processing and heat treatment condition. Consequently, they are essentially

non-magnetisible alloys even though they may be weakly attracted by a magnet

under certain conditions.

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9) New Cryogenic Materials: Fe-Mn-Al Alloys by J. Charles et. al, , ref. (15)

This latest paper on this subject has investigated the mechanical properties of these alloys at liquid nitrogen (-196°C or -320°F) temperature as well

as at room temperature. The base alloy studied was Fe-30Mn-5Al with increasing

amounts of carbon and silicon. They observed a high degree of dendritic

segregation or coring in the small ingots weighing only 2.6 lbs. Special

thermomechanical treatments were used for mechanical working of these small

ingots which involved 33% cold reduction, homogenization at 1150°C for 48 hours,

another 33% cold reduction, recrystallization at 925°C for 25 minutes. An ASTM

grain size of 4 was achieved ’ this way. The effect of carbon on the mechanical

properties at both room temperature and -196°C are shown in Fig. 9. The low

temperature properties are excellent with impact strengths of more than 32 ftIb./in^. Another set of alloys containing Fe-30Mn-5Al-0. 3C with increasing

amounts of Si showed the important hardening effect of Si at both room temperature and -196°C, Fig. 10, with no loss in ductility or toughness. The exceptionally high low temperature ductility and toughness of these Fe-Mn-Al

steels down to -196°C make them very attractive candidates for cryogenic

applications such as liquified natural gas containers, pipelines etc.

10) Mn-Al Steel Developed by Ford Motor Co., (ref. 16)

A very brief mention of this work exists in the 10th edition of Materials

Handbook. The steel composition was 30%Mn, 9%A1, l%Si, 1%C and had a tensile

strength of 120 ksi with 18% elongation. It work hardened rapidly and the coldrolled material responded to strain-aging to achieve a tensile strength of 300

ksi and yield strength of 290 ksi. The alloy was reported to have high resistance

to oxidation and sulfur gases up to 1400°F.

11) U.S. Seeking New Kinds of Stainless Steels - by P. Weintraub, ref. (17)

This report appeared in the American Metal Market/Metal Working News about

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one and a half years ago. Apparently there is some work sponsored by the U.S.

Bureau of Mines currently in progress at the Massachusetts Institute of

Technology to use the huge amounts of Mn nodules found at the ocean beds and

develop manganese-base alloys which might replace "chromium-laden" steels.

These alloys could contain up to 40%Mn, 10%A1, balance Fe. They could contain

some chromium for more corrosive environments. The report says "it's possible

that experiments will reveal the use of chrome has been more or less habitual

rather than necessary in a large number of cases". This point has been emphasized earlier. Yet, the feeling regarding Cr-substitution in many circles

is that of "waste of time and money" and "not do-able". The same feeling

prevails in the steel industry which has "no great impetus for activity" in

this vitally important area because they "see no imminent shortage of

chromium on the horizon". Personally, this attitude is simply appalling to

me. The article ends in a message which needs to be repeated over and over

again: "nobody wants to depend upon new and expensive technology . . . but

we would rather rely on new technology than a foreign supply of chrome".

Some Low-Chromium Substitutes

In addition to developing totally new chromium-free materials such as

those described above, the possibilities of reducing the chromium content of

existing stainless steels are also promising. NASA-Lewis Research Center

performed some work on reducing the chromium content of 304 stainless steel

which is the largest single stainless steel grade in use comprising over 30%

of all the stainless steel consumed in the U.S. (18). A minimum of 12% Cr was

found to be necessary to retain the corrosion and oxidation resistance comparable to 304SS. Some of the alloys explored in this study are shown in

Table 9. The last two austenitic alloys shown in this table exhibited

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oxidation resistance comparable to 304SS in high temperature (870°C or 1600°F)

cyclic oxidation tests, Fig. 11. Tests in copper sulfate-sulfuric acid solution show that the corrosion resistance of these alloys is also comparable to

the standard 304SS. To compensate for the reduced chromium content, 3.65%Si

in one of these alloys and 2% each of Al and Mo in the other were added. The

presence of these ferrite forming elements necessitated an increase in the Ni

content of these alloys to 10% and also an addition of 5.1%Mn in the Al-Mo

modified steel. The tensile properties of these alloys in the temperature

range of -321°to 1000°F were also comparable to the commercial 304SS. Although

these alloy compositions are not fully optimized yet, they clearly represent

the potential of 33% savings of chromium in currently used 304 stainless steels.

These low chromium modifications of 304SS could be used successfully in heat

exchangers, chemical process equipments, cryogenic containers, automobile

trims and a large variety of other applications where stainless steels are

currently used.

Another family of stainless steels based on the Fe-Mn-Al-Cr system have

apparently been developed on a laboratory scale at Fulmer Research Institute

Limited, U.K. (19). Typically these alloys could contain about 10-15%Mn,

4-5%Al, 8-ll%Cr, bal. Fe with yield stresses reaching about 600Mn/m^( 90ksi)

for 15%Mn alloys. They exhibit good oxidation and sulfidation resistance up

to 1000°C and are superior to 18/8 stainless steels. Figure 12 shows their

behavior in aggressive sulfide-bearing environments compared with the Cr-rich

oxidation resistant materials (20) . These alloys can be processed by production

forming routes similar to 304SS although higher forces are necessary for

equivalent deformations but no intermediate annealing is required. Some of

the industrial applications of these steels could include naptha reforming

P9-16
17

plants, gas and oil burners, automobile exhaust systems, desalination equipment etc

.

Review of Some Preliminary Work Spurred by the 1977-78 Papers

The publication of these papers (1,2) by the author on Cr-and Ni-free

stainless steels did attract quite a bit of attention. A large number of

inquiries were received concerning various properties and utility of the

Fe-Mn-Al alloys proposed in these papers. Since very little information was

available in the literature at that time and since I was not and I am still

not actively pursuing a research and development program on this subject,

most of those inquiries remained unanswered. Now, 3-A years later they are

still largely unanswered. However, I was fortunate enough to be in a position

to provide some samples of these alloys to those who were interested in doing

some exploratory work of their own. Some of them have now reported the preliminary findings of their work. Others, who were more stimulated by the idea

of the Fe-Mn-Al system made their own alloys for their investigations. Some

of them have also reported their early results. Some of these preliminary

results are quite encouraging and some are rather discouraging. Also, some

of these efforts are quite thorough and some are rather tentative. Therefore,

it should be remembered that some of the negative results may not necessarily

be a true reflection of the alloy but may simply be due to a lack of committment,

time, interest, funding, manpower etc. on the part of the investigator to

undertake a comprehensive and more systematic study. All these preliminary

results were reported to me privately and have not been published yet. Although it is a great pleasure for me to provide a brief sxjmmary of these early

results, it is not my intent to "let the cat out of the bag" or deprive the

original investigators of their recognition. It is only hoped that this brief

P9-17
18

exposure of their work will be beneficial in providing added stimulus to them

and others in further exploring the full potentials of these Fe-Mn-Al alloys.

- - - A joint program (21) was initiated between Foote Mineral Co. and Sandia

Laboratories to investigate the hot working characteristics of Fe-Mn-Al alloys

with the eventual goal of using these materials for nuclear waste containment.

The chemical composition of the alloys for this initial study are shown in

Table 10. Ninety-five pound ingots were melted under argon atmosphere using a

combination of high-carbon ferromanganese and electrolytic manganese and stock

aluminum. The examinations for macro segregation in the ingots showed that the

chemical compositions were remarkably homogeneous from one end of the ingot

to the other as determined by qualitative emission spectroscopy. This was

confirmed by quantitative chemical analysis. Microsegregation was examined by

measuring the dendritic arm spacing as a function of distance from the as-cast

wall. Attempts were also made to identify the chemistry of the various as-cast

constituents using x-ray and microprobe analysis techniques. The dendritic

arm spacing measurements were grouped between 1-2 x 10 in. without any

evidence for microsegregation. The results of these macro and microsegregation are in contrast to those of Charles et. al. (15).

Microprobe examination indicated that all of the alloys were 2-phase

in the as-cast state. One of these phases, both appearing austenitic by x-ray

analysis, tended to Mn rich, the other Fe rich. Probe analysis indicated the

Al and C were uniformly distributed. Further, Si also was uniformly distributed,

except for the S4 alloy, where there was some tendency for the Si to be associated with the Mn rich phase.

Homogenization : Homogenization temperatures were established by holding samples

for 4h at temperatures between 900 and 1200°C followed by air cooling. Table

11 presents the final homogenization temperatures selected for further study.

P9-18
19

SI, S2, S4, and S6 all responded in a similar fashion, that is above 1050-1100°C

the alloy was single phase, all memory of the cast structure having been eliminated by the homogenization treatment. S5 was single phase at 1050°C, 2-phase

(Y + 5) at 1100°C and still retained the as-cast structure at 1000°C. Finally,

numerous attempts to achieve a single phase austenite in S3 were unsuccessful.

It was two-phase up to 1200°C. S3 was therefore dropped from the program.

Hot Ductility : The hot ductility of the as-cast and homogenized Fe-Mn-Al alloys

was evaluated using a Gleeble test at an effective strain rate of 20 sec

Although this test procedure is rather crude, it does allow one to examine the

influence of test temperature on the ductility and thereby select an ingot breakdown temperature. Figure 13a-e summarize the results. All of the alloys exhibit

better hot ductility following homogenization. Apparently, the low ductilities

of SI, S4, S5, and S6 in the as-cast condition at temperatures greater than

1100°C is associated with a tendency for the failure to follow the interface

between the Mn and Fe rich phases. The behavior of S2 is anamolous in that

regard and is not understood at this time.

Based on these results, the following ingot breakdown procedure is suggested.

1. All alloys to be homogenized per Table 11.

2. Ingot breakdown with starting temperatures

SI 900

S2 1025

S4 900

S5 950

S6 1025

These temperatures assume some adiabatic heating during deformation. The

actual ingot breakdown according to the above schedule and extruding of bars

will now be pursued.

- - - Contrary to all the evidence in the literature regarding the workability

of Fe-Mn-Al alloys, a domestic steel company had considerable difficulty in

P9-19
20

cold rolling an Fe-30Mn-8Al-lC alloy (22) . They attempted to cold roll a

h" thick hot-rolled piece and surface ground for another purpose. No

solution annealing or homogenization was performed prior to cold rolling

attempts. In this condition it could not be cold rolled. So, it was hot

rolled to 0.1" thickness from 2150°F but edge cracking was experienced. The

edges were trimmed and then the cold rolling of the 0.1" thick material down

to 0.05" could be performed successfully. It is believed that if the original

material had been solution annealed and quenched prior to attempting the cold

rolling operation it would have been successful. The fast cooling rate of the

0.1" thick material after the intermediate hot rolling may have simulated the

solution anneal and quenching treatment, thus, making the subsequent cold rolling down to 0.05" possible. As mentioned earlier the properties of this type

of alloys are very sensitive to heat treatment and state of prior mechanical

working which must be kept in mind during processing of these materials.

- - - The Defence Metallurgical Research Laboratory in India has undertaken some

work on the austenitic 32Mn-9. 6Al-lC-Fe alloys to characterize its mechanical

properties, deep drawability, oxidation and corrosion resistance (23). Structural

characterization is also being carried out to understand the mechanism and kinetics

of aging. The room temperature mechanical properties of hot worked solution

annealed and aged specimens showed the typical age-hardening peak at about 600°C

aging temperature in agreement with previous observations of other workers

(5-10) . This peak strength corresponds with a minimum in ductility which is

still quite high. The Erichsen cup test for deep drawability showed identical

cup depths for the Fe-Mn-Al alloys and 304 stainless steel for 1.45mm thick

sheets. In 0.63mm thick sheets these alloys showed higher cup depths than 304

although the load requirement for the Mn-Al steel in both cases was about 30%

higher.

P9-20
21

100 hour oxidation test data showed these steels to be marginally inferior

at 700°C, and substantially inferior at 900°C to 304SS. Corrosion rates of

half immersed samples of these alloys in 7.5% NaCl were approximately double

compared to 304SS. In 50% HCl the results were considerably poorer than 304.

However, no attempts to etch away the two-phase surface layer or any other

surface preparations were made yet as suggested by Batrakov et. al. (11).

The corrosion results after such surface treatment will be of interest. Compositional modifications are also being planned to improve the corrosion and

high temperature oxidation resistance.

The cast, homogenized, and aged samples showed the presence of the austenite

I phase only. The cold-rolled samples and samples cooled to liquid nitrogen

temperatures also showed the austenite phase only. In the as-cast condition

the alloy was magnetic but in the homogenized and solution treated condition

it was non-magnetic. Aging between 500-800°C caused the alloy to be magnetic

again. Currently, work is in progress to study the kinetics of aging at

various temperatures. Thorough investigation of the microstructural details

and correlation with the properties is also underway.

- - - A "Femnal 30'* alloy containing Fe-30Mn-8Al-0.8C also showed its oxidation

resistance being limited above about 600°C (24) . This alloy was also found to be

unable to resist oxidation/sulf idation attack by multi-component gas mixtures

(H 2 S/H2O/H2 ) at 1800°F. The alloy disintegrated after about 100 hours of testing. However, resistance at lower temperature has not been investigated. The

preliminary results of an ongoing work on resistance to sulfidation in a reducing atmosphere of H2+I vol.% H2 S show this alloy to be the most resistant

material presently available, being 10 times more resistant than the 300 series

stainless steels. This is believed to be due to the formation ‘of a "peg-like"

duplex scale, consisting of an outer layer of MnS with an inner layer of AI2

S

3 ,

P9-21
22

which is highly resistant to spallation.

Thile it is too early to speculate, such alloys do seem very promising

as sulfidation resistant materials in reducing atmospheres. At present this

Femnal 30 alloy is being tested in a pilot plant oil gasifier in the U. K.

- - - Some early results of corrosion tests in the as-rolled condition of an

"Anodinox T-OA" alloy (Fe-36.5 Mn-7.A A1-0.13C) were very poor in 10%H2S0^,

1%HC1, 65%HN02, and 20%CH3COOH (25). A polarization curve in 10%H2SO^ showed

no passivation region for this alloy for a current density up to 1 A/Cm^.

Again, it must be remembered that the alloy was tested in the as-rolled condition with no removal of the Al- impoverished t-wo-phase surface layer. It

is hoped that the fully austenitic homogenized alloy after the removal of the

scale and the two-phase surface layer will give improved results. Some modification of the composition are also planned.

- - - A small foundry in Italy produced some cast test bars of an alloy consisting of Fe, 35.A%Mn, 9.75%A1, 1.7%Si, 0.79%C. The corrosion resistance of

these test bars in chlorinated media were found to be very poor compared to

316SS. However, no further details of the test, the surface condition of the

test bars and their heat treatment are known (26)

.

- - - The preliminary results from a large oil company in the U.S.A. indicate

that these Fe-Mn-Al alloys have very good corrosion resistance in carbonic

acid media (27). The alloys tested were essentially Fe-30Mn-8Al-lC with or

without 1.5%Si. These alloys were about 15-20 times more corrosion resistant

in this medium compared to C-steels and were nearly equivalent to 9Cr- or

9Cr- Mo steels. The silicon-containing alloy was even more corrosion resistant.

Therefore, these alloys may have some promise in routined non-severe "sweet

environments" of oil field applications.

- - - A company in the business of composite materials performed some preliminary

P9-22
23

tests with these alloys to explore the feasibility of using these steels for

high temperature composite material matrices which could replace some Crcontaining high temperature superalloys (28). The composition of the alloys

tested were Fe, 30%Mn, 8-10%Al, 0.5-1.25%C, 0-1.5%Si. First, they were able

to produce 10-12 mils thick sheets of these alloys. These were subjected to

standard vacuum diffusion bonding process. The initial small scale test of

bonding tungsten fibers to these alloys was successful. The next planned

step was the bonding of SiC filament reinforcement. The results of that are

not known.

Summary

It is quite clear that there are some conflicting observations and experiences about the behavior of these Fe-Mn-Al alloys. The mechanical properties

and aging characteristics are known and consistent to a certain extent. Some

scattered and preliminary information on some of the other properties have

emerged over the last 3-4 years although the lack of systematic and concerted

efforts in many cases make some of this information rather tentative. In

addition, the data-base on many other important engineering properties such

as machinability , weldability, creep, fatigue etc. are totally non-existant

.

However, the available knowledge on these alloys, albeit meager, do point to

the future potentials of these materials at least in some specific applications.

It is clear that thorough and systematic studies of this new class of materials

are required. They represent one possibility of chromium-free substitutes,

perhaps in some selected non-critical applications. If the chromium situation

is serious enough, which I believe it is, we cannot afford to ignore or discount it without serious and concerted efforts to evaluate its potentials.

Acknowledgements

The author is grateful to A. Gray of the American Society of Metals for

P9-23
24

continued encouragement and Foote Mineral Company for continued support of

the ideas presented here.

P9-24
25

References

1) S. K. Banerji; "An Austenitic Stainless Steel Without Nickel and Chromium",

Metal Progress, p. 59, April (1978).

2) S. K. Banerji; "An Austenitic Stainless Steel Without Nickel and Chromium",

Trans. Indian Inst. Met., Vol. 30, N. 3, p. 186, June (1977).

3) "Contingency Plans for Chromium Utilization", Publication //NMAB-335,

National Academy of Sciences, Washington, D.C. 20A18, (1978).

4) G. A. Watson; "Chromium-Today and Tomomorrow", Iron and Steel Maker, p. 19,

May (1980).

5) G. L. Kayak; "Fe-Mn-Al Precipitation-Hardening Austenitic Alloys", Metal

Science and Heat Treatment, 2, p. 95, (1969) - in English, translated

from Metallov. i Term. Obra. Metallov., No. 2, p. 13, Feb. (1969).

6) M. F. Alekseenko, G. S. Krivonogov, L. G. Kozyreva, I. M. Kachanova,

and L. V. Arapova; Metal Science and Heat Treatment, 14, (3-4), p. 187,

(1972) - in English, translated from Metallov. i Term. Obra. Metallov.,

No. 3, p. 2, March (1972).

7) P. J. James; "Precipitation of the Carbide (FeMn) 3AlC in an Fe-Al Alloy",

J. Iron & Steel Inst., p. 54, Jan. (1969).

8) G. S .Krivonogov, M. F. Alekseyenko, and G. G. Solov'yeva; "Phase Transformation Kinetics in Steel 9G28Yu9MVB", Phys. Met. & Metallog., (4),

p. 86, (1975) - in English-translated from Fiz. Metal. Metalloved.,

39, No. 4, p. 775, (1975).

9) L. I. Lysak, M. F. Alekseyenko, A. G. Drachinskaya, N. A. Storchak, and

G. S. Krivonogov; "Structural and Phase Changes in Steel 9G28Yu9MVB During

Aging", Metallogizika, 59, No. 4, p. 29, (1975) - in Russian.

10) N. A. Storchak and A. G. Drachinskaya; "Nature of Age Toughening of FeMn-Al-C Alloys", Fizika Metallov. i Metalloved., 44, (2), p. 373, (1977) - in Russian, avail, in English as "Strengthening of Fe-Mn-Al-C Alloys During

Aging", Phys. of Met. & Metallog., 44, No. 2, p. 123, (1978).

11) V. P. Batrakov, M. F. Alekseenko, L. A. Filimonova, G. S. Krivonogov,

L. G. Kozyreva, V. G. Sapozhnikova , and V. A. Molotova; "State of the

Surface Layer and Corrosion Resistance of Steel 9G28Yu9MVB", Prot. Met.,

10, (5), p. 487, (1974) - in English - translated from Zashchita Metallov.,

10, (5), p. 533, Sept. -Oct. (1974).

12) R. E. Cairns, Jr. and J. L. Ham; "Aluminum-Manganese-Iron Alloys",

U.S. Patent No. 3,111,405, Nov. 19 (1963).

13) J. L. Ham and R. E. Cairns, Jr.; "Manganese Joins Aluminum to Give Strong

Stainless", Product Engineering, 29, (52), p. 50, 22 Dec. (1958).

P9-25
26

lA) W. H. Richardson; "Iron-Aluminum Alloys", U.S. Patent No. 3,193,384,

July 6 (1965)

.

15) J. Charles, A. Berghezan, A. Lutts, and P. L. Dancoisne; "New Cryogenic

Materials: Fe-Mn-Al Alloys", Metal Progress, p. 71, May (1981).^

16) "Manganese-Aluminum Steel"; Materials Handbook, G. S. Brady, 10th Edition,

p. 497, McGraw-Hill.

17) P. Weintraub; "U.S. Seeking New Kinds of Stainless Stee!(s", American

Metal Market/Metal Working News, p. 8,24 Dec. (1979).

18) B. H. LeCerf; "Chromium Substitutes Worth A Second Look", Iron Age,

p. 42, 29 May (1978).

19) "New Family of Manganese Stainless Steels" - in Mn News, published by

Manganese Centre, 17 Avenue Hoche, 75008 Paris, France

20) P. L. Dancoisne; "Research and Development of Manganese", Materiaux & Techniques, Special Issue - Manganese, p. 127, Dec (1977), Paris, France.

21) H. J. Rack and S. K. Banerji, unpublished work, (1980).

22) J. C. Jasper, private communication, (1979).

23) M. L. Bhatia, R. Sivakumar, and Narasimha Rao, private communication,

(1981).

24) P. Tomas, private communication, (1981).

25) F. J. M. Boratto, private communication, (1980).

26) R. Boscarelli, private communication, (1980).

27) M. T. Orillion, private communication, (1981).

28) J. F. Dolowy, Jr., private communication, (1979).

P9-26
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P9-27
TABLE 2

U.S.

Chromium

Imports

from

Various

Sources,

1963-1976

(1,000

short

tons)

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P9-28

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TABLE 3 Prospects for Chromium Savings by Materials Substitution

and Redesign (1,000 short tons) - ref. (3)

Consumption

Forecast

for 1977

Potential Chromium :Savings

End Use

a Category 1 — b Category 2 —

Metallurgical

Wrought stainless and

heat-resisting steels 232 69.6 (30%) - 69.6 ( 30%)

Tool steels 6 0 (0%) 0 (0%)

Wrought alloy steels 43 6.5 (15%) 30 (70%)

Cast alloy steels 14 1.4 (10%) 8.4 (60%)

Alloy cast irons 8 0.8 (10%) 4.8 (60%)

Nonferrous alloys 14 0 (0%) 0.7 (5%)

Other 6 1.8 ( 10%) 1.8 (50J'o)

Subtotal

Refractories

323 80.1 (24.8%) 115.3 (35.7%)

Chrome and chromemagnesite brick 10 7.5 (75%) 2.5 (25%)

Magnesite-chrome brick

Granular chrome-bearing

14 4.0 (28%) 8.0 (56%)

materials 27 5.0 ( 18 . 5%) 20.0 (74%)

Granular chromite 10 4.0 (40%) 4.0 (40%)

Subtotal 61 20.5 (33.6%) 34.5 (56.6%)

Chemicals

Pigments 18 4.5 (25%) 9.0 (50%)

Metal finishing 15 7.5 (50%) 2.3 (15%)

Leather tanning 11 0 (0%) 11.0 (100%)

Drilling muds 3 0 (0%) 0 (0%)

Wood treatment 4 2.0 (50%) 2.0 (50%)

Water treatment 4 0 (0%) 2.0 (50%)

Chemical manufacture 5 0 (0%) 2.5 (50%)

Text! les 2 1.0 (50%) 1.0 (50%)

Catalysts < 1 0 (0%) 0 (0%)

Other 5 1.0 (20%) 3.0 (60%)

Subtotal 68 16.0 (2 3.5%) 32.8 (48.2%)

Total 452 116.6 (25.8%)^ 182.6 (40.3%)J

a— Technically feasible now. b“ Potentially feasible after 10 years of research and development,

c— Number in parentheses indicates percent saving of consumption

forecast for each end use category.

d— Percent saving of total for all categories.

P9-29
TABLE 4

Example

of

Possible

Chromium

Savings

in

Stainless

Steel

(tons) nou

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03 tj

03

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P9-30

The

figures

in

parentheses

indicate

the

percentage

of

possible

chromium

savings

in

that

category.
TABLE 5. Permeability of Fe-Mn-Al

Alloys after Quenching and after

Quenching + Aging 20 h at 600°C

- ref. (5)

Composition, % Permeability, G/Oc

quenched

C Mn A1 quenched + aged 20 h

at. 600°C

0.93 23.7 8.2 1,008 1,900

0.96 28.2 8.3 1,009 1,046

0.95 30.1 7.9 1,007 1,040

0.97 30.5 10.6 1,038 3,860

0.96 34.1 10.1 1,017 2,450

P9-31
Table 6: Composition and Mechanical Properties of some Fe-Mn-Al Alloys - ref. (12)

.4 II 0 V No. I’orccnl'Pcrccnt Percent Percent Percent i Test Heat Tensile Yield Percent Percent

A1 c Mn Other -Austenite

1 Temp.' Treat.! Sglh. Sgtli. E R .6.

133 IG 4 0.02 Mo-2.5...

i f> RT 110.000 10 0

1

{> 1.200 IRT

50.000 20.000 55 90

131 14 5 0.04 Ti—2.9.... 0! 2,000 Q

1.450Q

34,000 33.000 0

14 0

11,200 47,250 45.000 46 [RT I2.000Q 92,500 79.500 2 5 130 g.4 0.05 20.6

1 10 1 U.450Q

/2.000Q

107. 750

33, 750 78.000

33, 750 30 60

1 0 0

1 11.200 U,450Q 17.000 16.6at 187 99

172 15.3 1.08 28.3 20 [RT 2.000Q 2.000Q

168,000 74.250

132,000

12,000

4.5

166

2.4 U.200 98

j RT (2.000Q 141.000 14 14 173 12.2 0.67 25 2 40 >2,OOOFC...

ll,450Q

111.000 130.500 96,000

69. 45U 2.5 3.5

3

1 4 I1.2OO 2.000Q 55.750 45,000 22 24 IRT /2.000Q l2,000FC...

71. 750 28,500 55 n

17« 7.0 0.03 39.3 100 J 72.000 22»600 57 72

I1.2OO 2.000Q 41.250 18.000 88 66 IRT /1,450Q 75.000 69.000 0 0

190 10. 5 0.55 10.4 .Mo—2.8... 0 U.OOOFC...

1,450Q

64.900 45.000 0. 5 0.3

[1.200 IRT

62.500 12-000 62 65 /2.000Q lie. 260 10.. 750 10 2.0

210 14 6 0.02 MM-0.22 0

}I1.200 I2.OOOFC...

2,0000. .... 110.000 29,500

1. 5 1.0 21,000 92 87 (RT I2W0Q

12,000FC...

2,0000

109.000 55.000 73 71 213 10.2 0. 76 34. 4 100

j 110, 760 81 000 72 82

[l,200 60.500 40,500 34 32

1 |RT /2.000Q IIAOOO 61,000 62 73 214 - 7. 5 0. 7B 37 0 100 1 } 12.000FC... 116,760 64.750 66 71

jl.200 2,OOOQ 63, 760 33,000 66 64 |RT /2,OOOQ 119,000 65.500 45 51 226 10.0 0.27 33.8 Ni—2.3.... 60 J \2,OOOFC.-. 2,0000

97, 780 38,600

49.600 27.60025 30

ll,200 61 41 IRT /2,000Q 118,500 63, 750 45 52

227 12.5 0.30 33.0 .Nl—4.3.... 70 \2,000FC — naeco

40,500

55,500 30,000 35 32

I1.2OO 2^6060..... 37 31 IRT

(JCoOOFC—

123.000 63,250 -43 89 228 0.3 0.34 35.4 Ni—fi 4 80 J 114.000 49.000

62,500 35 32

[i.aoo 2,0000 %,000 31 81

> BT-Room Temperatorv; 1200-1200“ ?.

< The Dumber is tbe tempcrMure in de^ees F. that tbe alloy was held at prior to cooUd£. The describee the mauDer Ui which the alloy was cooled; Q-bot oil quench; FC -furnace cool.

• 8beet Sp^men NAVORD Report 4237 "AppbcaUons, Properties and Fabrication of Thermenol Type Alloys.”

\ P9-32
Table 7. Oxidation Test Results on Mn-Al

Steels of Table 6 - ref. (12)

(Depth of oxide penetration in inches)

Alloy No.

1,500'^F. - 100 hours

2,200^F. - 46.5 hours

133. .0005 .0005

134 .0005 .0005

136 .0005 .0015

172 .0005 .0005

173 .0005 .002

196 .0005 .001

210 .0005 .0005

213 .0065 .014

214 .0155 .047

226 .001 .0075

227 .002 .0035

228 .004 .005

310 stainless steel .0005 .0025

S-816 (cobalt base) .0005 .0625

U-500 (nickel base) .0005 .0025

Table 8. Composition of Mn-Al Steels

used in ref. (14)

Kxtmpio Al Mo C Cb Nl Otter

1 0. 46 21 1 1 04 0 30

J 10. 1 23 6 1.0 0.V2

3 10 61 24. 5 1. 1 ao3

i 10. V5 23. 6 1. 12 0.77

ft W 32 21.6 aon 0 84

ft V. 12 25 0 O.VK .0 oo

7 0. W5 25.0 1.04 0.70

H 0 04 24 7 1 01 3J7 W.

V 6.V3 21 0 a86 6.64 W.

10 0.80 24.5 1. 12 0.20 0.62 W.

11. . 0 23 28. 7 1.0 6.4 W.

12 V. 32 22.0 a 02 11 9 W.

13 10. 52 25. 5 1. 16 4.60 W.

14 10 30 26. 7 0. 76 0.60 Mo.

15 O.KO 24.0 0 04 0 81 S.26 Mo.

16 8.00 24.7 1.0 Nil 9.S8

17 8. 4 27.8 1. 0 14.98

18 8.63 27.2 1.00 0.36 ».6S

19 8.21 27.1 0 44 Nil Il.t /4.9 Cr.

\0 48 Tl.

20 10. 10 24.4 0.63 Nil 3.8 Cr.

21 8. 00 24. 8 0 OH 0 63 10 18 Cr.

22 7. 76 24.2 0.64 16.9 Cr.

23 10. 10 23 1 1 04 6 C4 2.10 Co.

24 0. 14 24. 4 0.06 Nil 1.66 Co.

25 0 G1 25.8 0. H6 2.1 Co.

28 0. 40 23 6 0 08 0. 70 0J6 Vo.

27 0.04 27.4 0.92 Nil 0.71 B. w 8. 00 24.4 0.82 Nil 0.92 TI. 0.02 24, 7 0. 86 0 60 0.64 Cu.

30 0. 34 24.0 0. V5 0. 74 2.24 Cu.

31 10.37 24.4 0. 82 Nil 2.0 Zt. 32 10. 78 24.8 1. 18 1.2 33 10.37 24.8 1. '20 Nil 0.4 Co.

34 0.48 26.1 0.90 0. 70

P9-33
Table 9. (ref. (18)

Composition of Modified 304 Stainless Steel

Alloys.

Ferritic—1st Generation, Austenitic— 2nd Generation

Alloy Nominal Chemical Composition, wt. %

Cr Ni Mn Si Al Mo C Fe

304 Stainless Steel 18 8 1.1 0.65 — — 0.06 Bai. 12% Cr-Si (Ferritic) 12 8 1.1 2.65 — — 0.06 Bal.

12%) Cr-AI-Mo (Ferritic) 12 8 1.1 0.65 2 2 0.06 Bal.

12%) Cr-Si (^Austenitic) 12 10 1.1 3.65 — — 0.06 Bal.

12%) Cr-Al-Mo (Austenitic) 12 10 5.1 1.65 2 2 0.06 Bal.

P9-34
Table 10. Chemical Composition of Foote-Sandia Alloys -ref. (21)

Alloy % C % Mn 1

—

< 1 % Si

S-1 0.93 29.6 7.98 0.09

S-2 0.57 29.6 8.03 0.05

S-3 0.17 29.7 7.77 0.01

S-4 0.99 29.7 7.71 1.49

S-5 0.59 29.2 7.59 1.52

S-6 1.00 29.6 7.87 0.73

Table 11. Homogenization Temperatures

Table 10.

for Fe-Mn-Al Ingots* of

Alloy Temperature (°C)

SI 1050

S2 1050

S3 —

S4 1050

S5 1050

S6

• 1100

Homogenization time

4h air cool

.

P9-35
ID Lf5 m ID LO If) (D m ^

NOIldlAinSNOO l^vinilAIObHO s n nv±01 dO iN30H3d

P9-36

FIGURE

1

Historical

Trends

in

the

Relative

Importance

of

Chromium

Usage

in

Metallurgical,

Chemical,

and

Refractory

Applications

(based

on

data

from

the

U.S.

Bureau

of

Mines

1950-1974,

1975,

1976)

-

ref.

(3)
<yb- kg/mm’; 6,

Aging Temp.

Fig. 2. Mechanical properties

of an alloy with 27.5% Mn,

9.1% Al, and 0.92% C as a function of aging temperature

(16 h) . - ref. (5)

I %x i-gi

lo*

• 9

"TT7 IN

S jnj50¥»999soos3oa»r ^

Aging tempernoge

5 tt O 20 %

Amoont of hatdeoktg phase

e \——

r -i-- i. T-j S a ts 20 7S JO Amouot of hardening phase

Fig. 3a Fig. 3b Fig. 3c

Fig. 3a) Variation of mechanical properties and amount of hardening

phase with aging temperature (8 h) for steel 9G28Yu9MVB.

Fig. 3b) Variation of mechanical properties with amount of hardening

phase for steel 9G28Yu9MVB.

Fig. 3c) Variation of toughness with amount of hardening phase in steel

9G28Yu9MVB:

1) grain size grade 4, 0.85% C; 2) grade 4, 0.95% C; 3) grade 1, 0.85% C.

- ref. (6)

P9-37
iOO 200 300 4 00- 500 600 700 SCO ACINC TEMPERATURE . °C

(a) (b)

Fig. 4 Hardness versus a) aging temperature (4 hrs.) and b) aging time

(at 600°C) for solution treated (1100°C, WQ) and as-rolled

Fe-9Al-9. 5Mn-. 96C-6. 77Ni alloy - ref. (7).

Fig. 5. Fe,Mn, and Al contents

vs depth of surface layer of a coldworked specimen of steel 9G28Yu9MVB

after heating in a furnace atmosphere

at 1100°C for 60 min. - ref. (11)

Fig. 6. Kinetics of growth of

two-phase surface layer of steel

9G28Yu9MVB at (°C) : 1) 1050;

2) 1100; 3) 1150. - ref. (11)

P9-38
P9-39
Fig. 8. Strain-Age-Hardening characteristics of coldrolled Fe-10.2Al-34.4Mn-0.76C alloy (#213)

after different amounts of cold reduction - ref. (12)

MPQ (KPpti) % MPq (lO^psi} %

Fig. 9. The effect of carbon on the mechanical properties of Fe-30Mn-5Al alloys.

Note that the properties at room temperature (A) are lower than those at liquid

nitrogen temperature (B) . TS=tensile strength; E^^and Eq strength (1 and

0.2% offsets); RA=reduction in area; A=elongation. -rer . (15)

P9-40
MPa(io^psi) % MPano^p'si)

Fig. 10. The hardening effect of silicon in Mn-Al steels is shown at room temperature (A) and liquid nitrogen temperature (B) . Again, better

results are obtained at cryogenic temperatures. TS=tensile strength;

El and Eo. 2=yi^ld strength(l and 0.2 offsets); RA=reduction in area;

A=elongation. -ref. (15)

P9-41
12 Cr-8 Ni-2 65 Si (Ferritic)

Time, Hr.

10

-50 200

(NPU\a

E

a;O)c

05 JZu

JZ

C)

05

o

a;a

c/3

CME

o

C55 E

0505C

05 O

05

05

O

05Q.

Fig. 11. Specific weight change data for modified 30A SS

alloys. Cyclic oxidation temperature, 870°C.

-ref. (18)

P9-42
Gain

de

polds

mg/cm

Weight

gain

mglcm^

Fig* 12. Beneficial effect of aluminium and manganese

on the sulphidation behaviour of stainless

steels at 900°C. -ref. (20)

P9-43
TEMPERATURE (K)

Fig. 13(a), Hot ductility test results for

ALLOY S-1 - REF. (21)

P9-44
Fig. 13(b). Same as (a) for alloy S-2 - ref. (21)

P9-45
#

Fig. 13(c). Same as (a) for alloy S-^ - ref. (21)

P9-46
Fig. 13(d). Same as (a) for alloy S-5 - ref. (21)

P9-47
Fig. 13(e). Same as Ca) for alloy S-6 - ref. (21)

P9-48
SUBSTITUTE STAINLESS STEELS WITH LESS CHROMIUM

Joseph A. Douthett

Armco Inc.

4

PIO

SUBSTITUTE STAINLESS STEELS WITH LESS CHROMIUM

BY

JOSEPH A. DOUTHETT

SENIOR RESEARCH ENGINEER

ARMCO INC.

TO BE PRESENTED AT THE

WORKSHOP ON CONSERVATION AND

SUBSTITUTION TECHNOLOGY FOR

CRITICAL MATERIALS

JUNE 15-17, 1981

VANDERBILT UNIV. - NASHVILLE

*'P^^^EiTl^fl5NS^
Steel analysts predict the stainless industry to be one of the

FASTER GROWING SEGMENTS OF THE AMERICAN STEEL MARKET. As A

STAINLESS PRODUCER^ ^ WE AT ArMCO ARE KEENLY AWARE OF THE

DEPENDENCE WE AS A COMPANY AS WELL AS A NATION HAVE ON FOREIGN

SOURCES OF SUPPLY FOR OUR CHROMIUM. We HAVE SEEN THE CHARTS ^

SHOWING OUR GREATER THAN 90% RELIANCE ON COUNTRIES LIKE SoUTH

Africa^ Russia and Zimbabwe. We view with alarm^ ^ projected

U.S. demands for Cr which will more than double by the year 2000.

Our predicament is further emphasized by comparing the U.S. vs.

THE World demand for Cr again through 2000 With no reserves^

WE WILL STILL ATTEMPT TO CONSUME OVER 20% OF THE PREDICTED WORLD

NEED. With some metals^ a secondary supply source comes in the

WAY OF RECYCLED SCRAP. ^ WiTH Cr ONLY XO % OF OUR U.S. SUPPLY

CAN BE GAINED VIA THIS ROUTE.

Being a Cr "have-not" nation^ we naturally look to who controls

THE AVAILABLE SUPPLY. A TABLE ^ OF LEADING PRODUCERS POINTS

OUT’ THE RATHER PRECARIOUS LOCATIONS OF RICH DEPOSITS. OUR

NEEDED Cr is CONTROLLED BY "Un" COUNTRIES: ^

Unstable

Unfriendly

Underdeveloped

Such countries are prime candidates to cartel the world's

CHROMIUM. Russia is presently actively promoting such a plan

TO Turkey and Zimbabwe.

PlO-2
The stage for such a cartel is even further set by reviewing

STEEL ALLOYING ELEMENT PRICE TRENDS FOR THE LAST 5 YEARS. ^

Of the common elements^ all have outstripped Cr. Like 30<I^/gallon

GASOLINE^ LOW PRICED READILY AVAILABLE Cr DESTROYS THE INCENTIVE

TO DEVELOP ALTERNATIVES. ThE ECONOMIC IMPORTANCE OF Cr IS

FURTHER EMPHASIZED ^BY REVIEWING THE COST/lB. OF COMMON STEEL

MELT ADDITIONS. FoR WHAT IT DOES FROM A CORROSION AND OXIDATION

RESISTANCE STANDPOINT^ Cr IS THE MOST COST EFFECTIVE STEEL ALLOYING

ELEMENT AVAILABLE. As BUREAU OF f^INES EXPERT JoHN MoRNING^^

PUT iT^ "There is no known substitute for chromium in stainless

steels". To replace Cr it would most likely cost the consumer money.

With no impending cutoff^ customers will not pay more simply to

CONSERVE Cr, even if the properties of a replacement alloy are

EQUIVALENT.

Due to economics and lack of an available substitute, American

INDUSTRY HAS BEEN CONDUCTING PROGRAMS OF PARTIAL REPLACEMENT

FOR Cr. Let's review some of these industry approaches and

THEN CHECK SEVERAL ArMCO IDEAS.

NASA HAS EXAMINED PARTIAL Cr REPLACEMENT WITH Mn, Si, Al,

AND Mo. While 33% of the Cr can be saved, the alloys have definite

PRODUCTION AND ECONOMIC PROBLEMS. HoWEVER, SUCH ALLOYS DUE PERMIT

SHIFTING THE STRATEGIC METAL PROBLEM FROM Cr TO Al, Mn, Mo, OR Nl

.

R. A. Lula of Allegheny Ludlum revealed a series of Cr conserving

AuSTENITICS 1^ THAT ALLOWED 35% OF THE Cr IN T30A TO BE REPLACED

BY Mn and N. Economically, a promising trade off although

QUESTIONABLE FROM A CORROSION STANDPOINT.

PlO-3
Nr. Lula further stated in his 1979 "Netal Progress" article

THAT STAINLESS CLAD HAD POTENTIAL. NOTE HERE THE Cr AND Nl

SAVINGS THAT T304 CLAD ONTO 409 WOULD GIVE. A 409 CORE ALSO

GREATLY ELIMINATES EDGE CORROSION SEEN WITH MOST COATED OR CLAD

STEELS. Cladding economics still remains a problem as would

WELDING.

Partial Cr replacements have been and are presenlty being studied

AT Armco. a great deal of effort has been expended on Cr

CONSERVING ALLOYS FOR THE AUTOMOTIVE INDUSTRY. As A STAINLESS

PRODUCER, Detroit is an exciting but unnerving place to visit.

Let me show you.

Stainless steel first found an automotive home in this the

CATALYTIC CONVERTER. 15 JuST ABOUT ANY CAR MADE IN THE 70'S HAS

ONE. There may be from 8-20# of stainless predominately 12%Cr

T409 IN YOUR CONVERTER. MULTIPLY ^THIS BY THE NUMBER OF VEHICLES

AND YOU HAVE A SINGLE INDUSTRY CONSUMING 11,000 TONS OF Cr

ANNUALLY.

However, it seems the converter was just the start of the

STAINLESS INVASION OF THE AUTOMOTIVE EXHAUST SYSTEM. NoW,

Detroit faces shaving weight to meet mandated fuel economies.

Small engines may run hotter. Here comes more potential stainless

APPLICATIONS. Sheet metal manifolds are predicted to go across

THE BOARD IN SEVERAL YEARS. ThESE MANIFOLDS MAY BE TUBULAR

OR STAMPED. 1^

PlO-4
The all stainless exhaust 20 is under study. Two car models at

G.fl. WILL GO TO A PREDOMINATELY 409 MUFFLER NEXT YEAR. NeDIUM

OR HEAVY DUTY GASOLINE TRUCKS 21 NOW OR WILL NEED CATALYTIC

CONVERTERS. PARTICULATE TRAPS 22 aRE BEING DESIGNED FOR DIESEL

ENGINES.

ThUS^ to make a long story SHORTy STAINLESS CONSUMPTION IN

Detroit could more than double within the next 5 years. Now

CARS ARE CRITICAL TO OUR DAILY LIVES^ BUT DO NOT REQUIRE THE

Cr as critically as do our aircraft engines or petroleum refineries.

Chromium may help improve our cars M.P.G. rating but what if

DUE TO A LACK OF Cr WE CAN't REFINE THE OIL TO MAKE THE GASOLINE.

As A STEEL CORPORATION^ WE ARE TORN. 23 Part OS US WANTS TO SELL

Detroit, the country's biggest steel user all the stainless it

WANTS. The other part of us looks downstream and urges Cr

CONSERVING ALTERNATIVES BE DEVELOPED. BoTH SCENARIOS ARE BEING

PURSUED.

When Detroit came looking for truck converter materials 24

MORE OXIDATION RESISTANT THAN 409, PAST STAINLESS TECHNOLOGY

SAID ADD MORE Cr. HoWEVER, BY USING Al IN SMALL QUANTITIES, OUR

OBJECTIVE WAS MET WHILE CONSERVING Cr. AdDING Al TO A STAINLESS

25

IS NOT AS EASILY DONE AS Cr7 WhILE Al DRAMATICALLY IMPROVES

OXIDATION RESISTANCE, IT IS DELETERIOUS TO OTHER ESSENTIAL

PROPERTIES. A BALANCE MUST BE STRUCK. ThE RESULT OF SEVERAL

YEARS OF RESEARCH IS A FAMILY OF "SR" CsCALE RESISTANT) ALLOYS.

PlO-5
These 2 materials ^6 known as 6 and 12SR are designed to replace

12 and 18Cr alloys respectively. The ,6%Cb addition dramatically

IMPROVES SHORT TIME CREEP OR SAG RESISTANCE. LIGHTER GAUGE

COMPONENTS CAN THEREBY BE DESIGNED.

Note this bar chart summary of GSR and 12SR cyclic oxidation

RESISTANCE. ThESE Cr-Al FERRITICS EFFECTIVELY FILL THE GAP

BETWEEN A09 AND 18Cr ALLOYS. As FOR CREEP ^8 qr SAG RESISTANCE

AT TEMPERATURES UP TO IGOOF^ THESE NEW ALLOYS FAR OUTDISTANCE

THEIR PREVIOUS FERRITIC COUNTERPARTS.

WHY DO WE EXPECT THESE ALJ.OYS TO GO AND NOT THE NASA Cr-AL“Si-F1o

ALLOYS REFERRED TO EARLIER? CoST EFFECTIVENESS. 29 JhESE ALLOYS

WILL GIVE COMPARATIVE PROPERTIES COMPARED TO MORE COSTLY AND

HIGHER Cr predecessors.

Besides truck converters and lighter gauge auto converters^

here's a diesel APPLICATION WHERE THE 12%Cr-Al ALLOY REPLACED

AN 18%Cr-10%Ni austenitic v;ith no thickness change.

The above alloys will conserve not replace Cr. An alloy

IDEA WHICH COMES MUCH CLOSER IS NP454 (NP STANDING FOR

NEW product). NP454 is an aluminized hot dip STEEL WHERE THE

SUBSTRATE INSTEAD OF BEING A LOW C STEEL IS A LOW ALLOY VERSION.

Adding Cr^ Al> Ti to the base metal increases the coated product's

UPPER OXIDATION LIMIT TO 1500F VERSUS 1250F. ^2 $UCH AN ALLOY

COULD REPLACE 409 IN SOME EXHAUST APPLICATIONS AT A LOWER COST

AND A SIGNIFICANT SAVINGS IN Cr . ALUMINUM COATING LOW ALLOY

AND STAINLESS BASE METALS MAY BE AN ANSWER TO OXIDATION PROBLEMS

PlO-6
WITHOUT INVESTING IN NEW CAPITAL EQUIPMENT.

While adding Al and Si to stainlesses^ how about an austenitic.

IN1212 IS AN International Nickel alloy designed jo conserve Cr

AND REPLACE T309 OR T310 STAINLESS. OXIDATION TESTS SHOW

THIS MATERIAL A PROMISING HIGH TEMPERATURE CANDIDATE.

While conserving Cr in T304^ let's try a different approach

BY OFFERING A COMPLETELY REDESIGNED COMMODITY AUSTENITIC STAINLESS.

NP599 IS A 16%Crv 8%Nn^ 2.25%Ni austenitic which uses Mn, N and

Cu TO REPLACE COSTLY Nl WHILE QUIETLY CONSERVING 2%Cr . WhILE

NOT QUITE AS CORROSION RESISTANT AS T304 NP599 OFFERS

BETTER STRENGTH AND A COST INDUCEMENT. ThE OBJECTIVE OF NP599

IS NOT TO OFFER A LOWER CORROSION RESISTANT T304 BUT TO PROVIDE

A COST EFFECTIVE MATERIAL ALTERNATIVE TO APPLICATIONS WHERE T304

IS ALREADY AN OVER-DESIGN. Cr CONSERVATION IS AN UNWRITTEN

SIDE BENEFIT.

How ABOUT A CORROSION HEAVY WEIGHT? ArMCO'S NITRONIC 50

EFFECTIVELY COMPETES FROM A CORROSION AND PRICE STANDPOINT WITH

MUCH HIGHER Nl^ Mo AUSTENITICS BY USING N TO AID IN PITTING

RESISTANCE. REPLACING Mo WITH N ALSO LIMITS THE AMOUNT OF Nl

NEEDED AS N IS A POWERFUL AUSTENITE STABILIZER,

Finally. I give you NITRONIC 60. a 4% Si'austenitic developed

FOR GALLING AND WEAR TESTS. TeSTS LIKE THESE INDICATE THAT

SOME APPLICATIONS MAY ALLOW US TO USE AN Fe BASE STAINLESS FOR

HIGH Co MATERIALS.

PlO-7
This pretty much summarizes the kind of material conservation

RESEARCH BEING DONE AT ArMCO . PERHAPS IT'S APPARENT THAT IN

A TIME OF READILY AVAILABLE Cr^ THE ONLY CONSERVING THAT CAN BE

DONE IS A PARTIAL REPLACEMENT IN STAINLESS ALLOYS OR THE DEVELOPMENT OF LOW Cr coated MATERIALS, HoWEVER^ FOR THESE TO BE

EFFECTIVE^ A COST ADVANTAGE INDUCEMENT MUST BE BUILT-IN. WHILE

MATERIAL INERTIA FIGHTS US IN EXISTING MARKETS^ LET's HOPE

EMERGING HIGH ALLOY MARKETS CAN BE GUIDED TOWARD STRATEGIC MATERIAL

CONSERVING ALLOYS.

PlO-8
PlO-9
U.S. IMPORT RELIANCE FOR SOME IMPORTANT METALS

Al^ Crv Co^ Mn^ Ni^ Nb> Ptj T/

Fe. V. Ti.

Si^ Cu^ Rare Earth Elements

Mo. Mg

PIO-IO
PROJECTED U.S. DEMAND FOR CR

YEAR DEMAND CIOOO TONS)

1976 461

1985 670

2000 1100

SOURCE; US BUREAU OF MINES
CUMULATIVE U.$. AND WORLD DEmD/RE$ERVE$ FOR CR

YEAR 1976-2000

(MILLION TONS)

U.S. WORLD

DEMAND 20 87

RESERVES <•1 830

SOURCE: U.S. BUREAU OF MINES
SCRAP AS A SECONDARY SOURCE

OF THE U.S. METAL SUPPLY

SOURCE: NETSCHERT, "JOURNAL OF METALS". MARCH 81

PlO-13
6

(MILLION i^iORT TONS)

LEADING CHROMIUM

PRODUCING NATIONS

SOUTH USSR ALBANIA ZIMBABWE TURKEY FINLAND PHILIPPINES

AFRICA

SOURCE: B. NETSCHERT, "JOURNAL OF METALS". MARCH 81

PlO-14
"UN"

UNSTABLE

UNFRIENDLY

UNDERDEVELOPED

PlO-15
8

NETAL PRICE TRENDS 1975-1980

CR

% INCREASE

112

m 115

cu 145

NI 146

FE 153

TI 290

MO 310

CO 625

PlO-16
PlO-17
PlO-18
PlO-19
PlO-^0
PlO-21
PlO-22
PlO-23
II

PlO-24
PlO-25
PlO-26
PlO-27
PlO-28
PlO-29
PlO-30
PlO-31
PlO-32
Al

in

Stainless

Steels

PlO-33

no

effect

negative

effect
PlO-34
PlO-35
PlO-36
PlO-37
PlO-38
PlO-39
PlO-40
PlO-41
f < K I 0^ w B CO

IB J>

0>

Qi CQ o o to • • CM

lo

q

CM 0>

to to • • q o

CM 00

to• o• CO o

T- CM

to o

CM to

CM CM o h- CO o o O CO

T“

O

CM

CM a> o

CJ T""

CO CO

PlO-42
PlO-43
TYPICAL CHEfllSTRY

NP599

COfTOITY AUSTENITIC

%L Wn Ei %Cr Mi ^Cu M

NP599 .035 8.0 .50 16.0 2.25 .75 .20

T304 .060 1.5 .50 18.3 8.50 .20 .03

PlO-44
NP599 vs. im

STRENGTH

CORROSION RESISTANCE

COST

CR CONSERVATION

NP599

1/

(//

DENOTES COMPARATIVE ADVANTAGE
PlC-46
TYPICAL

COMPOSITION O

<D

O

C

O

• o X o

CD o o • 00 S o LO• vP ai

1 1 o

o o o q

• o in

• cd

o CO T"

c

o

o

CO

E o

o

^

JZ •—

O z

PlO-47

Nitrogen

.08-.1

8
PI 0-48
PlO-49
SUBSTITUTION OF MODIFIED OCR-IMO STEEL

FOR AUSTENITIC STAINLESS STEELS

Vinod Sikka

Oak Ridge National Laboratory

pii
SUBSTITUTION OF MODIFIED 9 Cr-1 Mo STEEL

FOR AUSTENITIC STAINLESS STEEL*

V. K. Sikka

Paper to be presented at

Workshop on Conservation and Substitution

Technology for Critical Materials

Vanderbilt University

Nashville, Tennessee

June 15-17, 1981

*Research sponsored by the Office of Reactor Research and

Technology, U.S. Department of Energy, under Contract W-7405-eng-26

with the Union Carbide Corporation.

By acceptance of this article, the

publisher or recipient acknowledges

the U.S, Government's right to

retain a nonexclusive, royalty-free

license in and to any copyright

covering the article.

Pll-1
SUBSTITUTION OF MODIFIED 9 Cr-1 Mo STEEL

FOR AUSTENITIC STAINLESS STEEL*

V. K. Sikka

ABSTRACT

This paper describes the current program to develop a high

strength ferritic-martensitic steel. The alloy contains Fe—9% Cr—1% Mo

with small additions of V and Nb and is known as modified 9 Cr-1 Mo.

Its elevated-temperature properties and design allowable stresses match

those of type 304 stainless steel for temperatures up to 600°C and

exceed other ferritic steels by factors of 2 to 3. The improved

strength of this alloy will permit its use in place of stainless steels

for many applications. We expect this possible substitution to reduce

the demand for imported chromium.

INTRODUCTION

Chromium is an essential constituent in alloys used in oil

refining, petrochemical plants, conventional and nuclear power plants,

tanker trucks, gas turbines, industrial machinery, and all stainless

steel applications where high elevated-temperature strength and

corrosion resistance are required. Yet our major sources of supply are

South Africa and the Soviet Union, and in 1980 we imported 91% of the

chromium we required. The United States is currently following a policy

of stockpiling critical elements; however, the stockpile of chromium is

*Research sponsored by the Office of Reactor Research and

Technology, U.S. Department of Energy, under Contract W-7405-eng-26

with the Union Carbide Corporation.

Pll-2
2

180,000 tons short^ of a stockpile goal of 1.35 million tons. The

long-term solution to the critical element (Ft group, Cr, Co, Mn, and

Ta) shortage is increased metallurgical research that will lead to the

development of alloys that will reduce the shortage of these strategic

materials. The purpose of this paper is to present the status of

development of a modified 9 Cr-1 Mo alloy that can be substituted for

stainless steels in many applications and, therefore, will reduce the

consumption of chromium.

BACKGROUND

The 300 series austenitic stainless steels are the most common

structural materials for elevated temperature service. Typically, they

contain 18% Cr and 8% Ni, and they provide good air and steam oxidation

resistance. However, they do have problems with stress corrosion

cracking in media containing halides. Ferritic steels on the other hand

contain much lower chromium and nickel than to stainless steels and are

relatively free from stress-corrosion cracking. They do have less

oxidation resistance to air and steam. However, there are many

applications where ferritic steels of the 9 Cr-1 Mo type would be

acceptable if they had sufficient elevated-temperature strength. If

their strength were improved, ferritic steels would also provide better

resistance to thermal stresses because of their lower expansion

coefficient and higher thermal conductivity. Most of all, using an

alloy of this type instead of stainless steels would produce a

Pll-3
3

substantial saving of chromium. The following section describes the

properties of the modified 9 Cr-1 Mo alloy and compares them with those

of type 304 stainless steel.

RESULTS AND DISCUSSION

Modified 9 Cr-1 Mo steel is strengthened by small additions of

vanadium and niobium. The recommended ranges for various elements are

as follows:

c Mn Si P S 0.08-0.12 0.30-0.50 0.2-0.

5

0.02 max 0.01 max

Cr Ni Mo V Nb N

8-9 0.2 max 0.85-1.05 0.18-0.25 0.06-0.10 0.03-0.0

The alloy is currently recommended for use in the normalized and

tempered condition (1038®C for 1 h, air cool to room temperature, 760“C

for 1 h followed by air cool to room temperature). However, work in

progress shows that the desirable elevated-temperature strength may also

be obtained in the isothermally annealed condition (1038°C for 1 h, fast

cooled to 704°C, held for 24 h, then air cooled to room temperature).

The alloy has been commercially melted by air induction and

argon-oxygen deoxidation (AOD) and refined by the electroslag reraelting

(ESR) process. Ingots have been fabricated into tube, plate, and bar.

The fabrication processes used include hot forging, hot rolling, hot

extrusion, hot rotary piercing, centrifugal casting followed by cold

pilgering and cold drawing. No problems have been encountered during

any of these melting and fabrication procedures.

Pll-4
4

Specimens from commercial heats have been tensile, creep, Charpy

impact, and fatigue tested. Tensile data were used to specify the

room-temperature values of 414 MPa (60 ksi) and 586 MPa (85 ksi) for

yield and ultimate tensile strength, respectively. The request for

inclusion of specifications for this alloy in the ASTM Standards book is

currently under review by the appropriate committees.

The average and minimum curves defined by average values and

average values minus two standard deviations for yield and ultimate

tensile strength of standard and modified 9 Cr-1 Mo alloy are plotted as

a function of test temperature in Figs. 1 and 2. A minimum curve for

modified 9 Cr-1 Mo based on room-temperature specified values are also

plotted in these figures. The^ASME Code Case N-47 minimum value curve

for type 304 stainless steel is also shown for comparison. These curves

show the following:

1. The yield strength for both standard and modified 9 Cr-1 Mo

steel is significantly above type 304 stainless steel up to 700°C.

2. The ultimate tensile strength of both standard and modified

9 Cr-1 Mo steel is higher than that of type 304 stainless steel for

temperatures up to 550°C.

3. The modified alloy shows higher yield and ultimate tensile

strength than the standard alloy over the entire temperature range.

The standard 9 Cr-1 Mo data used in Figs. 1 and 2 come from the

United Kingdom (UK) and are on material in the normalized and tempered

condition. In the United States the annealed condition is typical for

standard 9 Cr-1 Mo. Data on modified 9 Cr-1 Mo steel are for the

normalized and tempered condition listed above.

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5

Figure 3 illustrates the variation in estimated average 10^ ^

rupture strength with temperature for various materials. Shown for

comparison are average values for type 304 stainless steel that were

used in the analysis^ to calculate the minimum values now given in

Code Case N-47, (This time was chosen because it is the longest time

for which verifiable strengths can be calculated for all data sets.)

The modified 9 Cr-1 Mo alloy is comparable in creep strength to type 304

stainless steel for temperatures up to about 625°C. Above 625°C it

begins to fall below the stainless steel in strength. The

12 Cr-1 Mo-V-W-0.2 C (HT9) maintains a strength comparable to stainless

steel only up to about 550°C before it drops off to lower values. The

2 1/4 Cr-1 Mo and annealed standard 9 Cr-1 Mo are lowest in strength,

and the UK normalized and tempered standard 9 Cr-1 Mo is intermediate.

Above about 675°C, however, all these ferritic steels except modified

9 Cr-1 Mo converge to about the same low strength value. Even as high

as 700°C the modified alloy remains intermediate between the stainless

steel and the other ferritics.

Fatigue data on modified and standard 9 Cr-1 Mo alloy are compared

with the average curve for type 304 stainless steel in Fig. 4. The

modified 9 Cr-1 Mo data were obtained in both air and vacuum. This

figure shows that the high-cycle fatigue properties of modified

9 Cr-1 Mo are an order of magnitude better than those of type 304

stainless steel. Because of its lower thermal expansion and higher

thermal conductivity, the advantage of modified 9 Cr-1 Mo alloy over

type 304 stainless steel under thermal fatigue conditions is expected to

be even better.

Pll-6
6

A very important consideration in the use of a given alloy is its

design allowable stresses {Sq), For a ferritic steel, Sq is given as

the lowest of the following four stress values at a given temperature;

1. Tensile; one-fourth of the tensile strength at temperature,

which is defined as the smaller of (a) the specified minimum tensile

strength at room temperature or (b) a value 10% greater than the minimum

tensile strength at temperature.

2. Yield; five-eighths of the minimum yield strength at

temperature.

3. Rupture; (a) 0.67 of average rupture stress for 100,000 h or

(b) 0.8 of minimum rupture stress for 100,000 h.

4. Secondary (or minimum) creep rate; the average secondary creep

rate for 0.01% creep per 1000 h (equivalent to 1%/100,000 h).

The estimated allowable stresses based on the criteria listed above

are plotted in Fig. 5. This figure includes values for 2 1/4 Cr-1 Mo,

standard 9 Cr-1 Mo, and type 304 stainless steel. It shows that

modified 9 Cr-1 Mo steel has higher allowables than the other ferritic

materials over the entire temperature range from 427 to 704°C. The

modified alloy is also estimated to have allowable stresses equal to or.

greater than those of type 304 stainless steel to about 600“C. For a

given design stress, modified 9 Cr-1 Mo alloy can be used about 75°C

higher than other ferritics and thus has an excellent chance of

replacing type 304 stainless steel.

To gain commercial experience with modified 9 Cr-1 Mo alloy, we

have installed tubes in two conventional utility power plants. At a

Tennessee Valley Authority power plant the modified tubes replaced

Pll-7
7

type 321 stainless steel tubes and have operated successfully since

May 1980. At a fossil plant of American Electric Power, the modified

alloy replaced type 304 stainless steel, and these tubes have been in

operation since April 1981.

Work is continuing in the areas of mechanical properties, physical

properties, weldability, and thermal aging effects. Initially we expect

to obtain approval of modified 9 Cr-1 Mo in Sections I and VIII (for

nonnuclear applications) of the ASME Code by 1982 or 1983. The nuclear

application of this alloy will require a substantially larger data base

and thus ASME Code approval for Section III is not expected until 1985.

SUMMARY

Although chromium is a critical element in many materials

applications, over 91% of the chromium used by the United States is

imported. To minimize our dependence on foreign sources of supply

requires alloys that use less chromium than do present alloys such as

stainless steels. This paper has outlined the development of a modified

9 Cr-1 Mo alloy, which uses half of the chromium and a trace amount of

nickel compared with stainless steel, but has elevated-temperature

properties matching those of type 304 stainless steel up to 600°C. The

design allowable stresses of this alloy are factors of 2 to 3 higher

than those of other ferritic steels in the temperature range of

500—700°C. For a given design stress, the modified alloy can be used at

75°C higher than other ferritic alloys that have from 2.2 to 12% Cr.

Improved elevated-temperature strength along with excellent physical

properties of this modified alloy are expected to make it quite

resistant to thermal fatigue.

Pll-8
8

Data on the modified 9 Cr-1 Mo alloy presented in this paper show

that there may be many applications where it could replace stainless

steel and thereby reduce the need for imported chromium.

REFERENCES

1, F, Warshofsky, "Strategic Minerals the Invisible War," Reader's

Digest

f

Feb. 1981, pp. 81—85.

2. R. A. Moen and L. D. Blackburn, Hanford Engineering Development

Laboratory, Private Communication, 1974.

Pll-9
1000

900

800

700

600

500

400

300

200

100

0

ORNL-DWG 81-4316

SHADED REGIONS REPRESENT ESTIMATED SPAN

FROM AVERAGE TO AVERAGE MINUS 2SEE

MOD 9 Cr- 1 Mo STEEL

STD 9 Cr - 1 Mo STEEL

100 200 300 400 500 600 700 800 900

TEST TEMPERATURE (°C)

Fig.

odif led

1. Comparison of 0.2% yield strength curves

9 Cr-1 Mo and type 304 stainless steel.

for standard and

Pll-10
1000

900

800

700

600

500

400

300

200

100

0

St

ORNL-DWG 81-4314

SHADED REGIONS REPRESENT ESTIMATED SPAN

FROM AVERAGE TO AVERAGE MINUS 2SEE

''Xy/x MOD 9 Cr1

Mo STEEL

STD 9 Cr - 1 Mo STEEL

100 200 300 400 500 600 700 800

TEST TEMPERATURE (°C)

Fig.

mdard

2. Comparison of ultimate tensile strength curves for

and modified 9 Cr-1 Mo and type 304 stainless steel.

900

Pll-11
500

450

400

350

300

250

200

150

100

50

0

A

Fi] 3 . Variation

materials.

of 10^-h creep-rupture strength with temperature

Pll-12
ORNL-DWG

79-12977

CO

LiJ CC

Z)

o

Ico

LU

_l

o

>

o

Pll-13

Fig.

4.

Comparison

of

continuous

cycling

fatigue

data

on

modified

9

Cr-1

Mo

steel

with

standard

9

Cr-1

Mo

steel

and

type

304

stainless

steel.
180

160

140

120

100

80

60

40

20

0

4

Sq>

0RNL-DW6 80-8020R

SOLID LINES FROM CURRENT ANALYSIS

OTHER RESULTS FROM ASME CODE

• TYPE 304 STAINLESS STEEL

O AMERICAN COMMERCIAL

9Cr-1 Mo STEEL (N & T FORGING)

0 AMERICAN COMMERCIAL 9 Cr - 1 Mo STEEL

Fig. 5. Variation of estimated design allowable stress intensity,

rith temperature for several materials.

Pllnl4
COBALT. A VITAL ELEMENT IN THE AIRCRAFT ENGINE INDUSTRY

Joseph R. Stephens

NASA-Lewis Research Center

P12
r
CONSERVING COBALT

Cobalt, A Vital Element In The Aircraft Engine Industry

Joseph R. Stephens

NASA Lewis Researcn Center

Cleveland, OH 44135

Cobalt Session Keynote Paper

for

Department of Commerce's

Workshop on Conservation and Substitution

Technology for Critical Materials

(June 15-17, 1981)

P12-1
ABSTRACT

Recent trends in the United States consumption of cobalt indicate that

superalloys for aircraft engine manufacture require increasing amounts of this

strategic element. Superalloys consume a lion's share of total U.S. cobalt

usage which was about 16 million pounds in 1980. In excess of 90 percent of

the cobalt used in this country was imported, principally from the African

countries of Zaire and Zambia. Early studies on the roles of cobalt as an

alloying element in high temperature alloys concentrated on the simple Ni-Cr

and Nimonic alloy series. The role of cobalt in current complex nickel-base

superalloys has not been well defined and indeed, the need for the high

concentration of cobalt in widely used nickel-base superalloys has not been

firmly established. This paper'will review the current cobalt situation as it

applies to superalloys and describe briefly the opportunities for research to

reduce the consumption of cobalt in tne aircraft engine industry—a field of

research being addressed by NASA's COSAM (Conservation of Strategic Aerospace

Materials) Program.

INTRODUCTION

Periodically the United States metallurgical industry is faced with

shortages of strategic materials that are critical to the economical health of

our country. In the early 1970's the potential reduction of chromium supply,

brought about in part by U.S. government sanctions against Rhodesia (now

P12-2
1 2 Zimbabwe), led to a number of studies , workshops , and research

3 4

efforts ’

. NASA Lewis Research Center's research efforts focused on

reducing chromium in 304 stainless steel which constitutes the largest single

use of chromium in this country. Results indicated that substituting aluminum

and molybdenum for a portion of the chromium produced an austenitic alloy with

comparable properties to that of 304 stainless steel and with the potential of

3

conserving one-third the chromium in this alloy .

% In the late 1970's another strategic metal shortage occurred in this

country which had a major impact on superalloy producers. Because of

political instabilities in Zaire, cobalt supply to the United States was

disrupted. Producer cobalt supply to alloy melters in 1978 and 1979 was put

on an allotment of 70 percent of 1977 usage. In addition to these

disruptions, a titanium sponge shortage plagued U.S. alloy producers in the

late 1970's.

Recent developments indicate that official policies of the United States

are now taking materials shortages seriously. The National Materials Policy

c. Act was passed into law during 1980. This Substitutioh Workshop, sponsored

by the Department of Commerce, is in direct response to this Act and a report

of this Workshop and other activities will be delivered to the Congress by the

Department of Commerce later this year. It is interesting to note that the

Departmeht of Commerce chose the aerospace industry as its case study of

strategic materials. We at NASA Lewis Research Center have formulated the

COSAM (Conservation of Strategic Aerospace Materials) Program^ with special

interest on the four strategic elements cobalt, columbium, tantalum, and

P12-3
chromium. Initial researcn efforts, underway since early 1980, are

concentrated on determining the role of cobalt in superalloys. Superalloys

constitute the single largest U. S. usage of cobalt. Therefore, the papers

presented in this two-part session on cobalt are very timely and of extreme

interest to the United States economy. An inoication of the importance placed

on cobalt by the current administration is the recent announcement that 3100

million would be used to build-up the U. S. stockpile^. Cobalt was singled

out as the first commodity to be purchased under this plan.

This paper will present the cobalt situation in the U.S. as it exists now,

describe the role of cobalt in nickel-base alloys, and describe the initial

COSAM Program on cobalt which was undertaken in 1980.

GEOGRAPHICAL CONSIDERATIONS

In planning and establishing our COSAM Program, strategic materials were

defined as those predominantly or wholly imported elements contained in the

metallic alloys used in aerospace components which are essential to the

strategic economical health of the U.S. aerospace industry. Cobalt was

identified as being within this definition of a strategic material. The

United States imports over 90 percent of the annual consumption of cobalt.

Figure 1 shows the leading producers of cobalt. The African countries of

Zaire and Zambia are noted to be the leading cobalt producers. Also listed in

g figure 1 are the sources of U.S. imported cobalt in 1980 . Zaire and Zambia

supply over 70 percent of the U.S. imports. This producer and supplier

P12-4
situation places the United States in an extremely vulnerable position if

supplies are cutoff for any length of time. An example of this occurred in

1978 and 1979 because of the political disruptions that Zaire experienced in

those years.

The objective of the COSAM Program is to provide technology options which

will support the aerospace industry in making strategic economic decisions

aimed at significantly reducing strategic metal (such as cobalt) consumption.

This Workshop along with other activities will help focus attention on our

non-fuel strategic material situation.

ECONOMICAL CONSIDERATIONS

Most of you are aware of the price increase in cobalt that has taken place

since 1977. Cobalt metal that was selling at arounu J5.50 per pound in 1977

increased to over $30.00 per pound during 1979 with spot prices as high as

$55.00 per pound. A historical rule-of-thumb has been that the price of

coDalt is typically higher than that of nickel by a factor of two to three

times. In 1980 that factor was in excess of seven times.

It is interesting to compare the price increase in cobalt with that of

oil, a commodity that most of us are concerned with. The average prices for

the year 1972 were chosen as a basis of comparison since this was just before

9

oil prices skyrocketed. The price of cobalt was $2.45 per pound while OPEC

oil sold for $2.28 per barrel^^. Figure 2 shows that OPEC oil experienced a

rapid rise in price during 1974 and again in 1979 and 1980. Interestingly

enough, cobalt price increase has kept pace with that of oil over the past

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eight years. Just as oil price increases have influenced the use pattern of

this commodity, the price of cobalt has brought about changes the use pattern

of this commodity as well. Some of the trends in cobalt usage will be

discussed in a latter section of this paper.

Even though the changes in prices of oil and cobalt have been similar over

the last eight years, the response by the American public and press has been

quite different. There has been a strong demand for conservation of and

substitution for oil, but the cobalt situation has received only minor

attention. The reason is the difference between the place of oil and cobalt

in the economy. A cutoff of oil for a few months could severly cripple our

economy particularly in the area of transportation and home heating. In

contrast the average consumer would not even know a cutoff of cobalt had

occurred unless it extended for a prolonged length of time.

A second reason is the tremendous demand for oil in this country compared

to most any other commodity. Annual consumption of oil is compared to that of

cobalt in figure 3 along with two other mineral commodities. Oil^^ is

g consumed by a factor of over five orders of magnitude compared to cobalt .

Iron ore, the largest used nonfuel mineral commodity, is overshadowed by oil

consumption. With the volume of oil imported (over one-half of consumption) a

significant effect on the bal ance-of-payments exists with a slight increase in

price. In contrast, with the relative small amount of cobalt imported (even

though over 90 percent is imported), a price increase comparable to that of

oil has a negligible effect on the bal ance-of-payments . These factors suggest

that the concentrated users of cobalt must take the lead in identifying

effective substitutes for cobalt that will permit the United States to remain

competitive in the world market.

P12-6
USAGE CONSIDERATIONS

Primarily because of the spiraling cost of cobalt, the United States has

experienced a decline in cobalt usage since 1978. Figure 4 shows that 20

million pounds of cobalt were consumed in 1978 and that in 1980 usage was down

to 16 million pounds, a 20 percent reduction in usage in only 2 years. During

this same time period the use of cobalt to produce superalloys, primarily for

aircraft engines, 'increased from 4 million pounds in 1978 to 7.2 million

pounds in 1980. This change represents an 80 percent increase in cobalt usage

in superalloys in only two years. The numbers are from the Department of

12

Interior's Mineral Industry Surveys

The importance of cobalt to superalloy production is illustrated in figure

5. It should be noted that 45 percent of the 16 million pounds of cobalt

consumed by the United States in 1980 went for superalloy production. So even

though total U.S. usage is down, superalloy usage continues to grow. The

shift away from cobalt usage has been achieved primarily by the magnetic

materials industry and by tne manufacturers of tool bits and dies. This nas

been accomplished by utilizing ferrite magnets and development of cobalt-free

cutting tools.

The increased usage of cobalt in superalloys can be attributed largely to

the increased orders of aircraft engines which is the predominate superalloy

market. Airlines have announced orders for new planes to supplement their

current fleets, and the new engines will require large amounts of the

strategic metal cobalt. Since the aircraft industry is a major factor on the

P12-7
13 positive side of the U.S. balance-of-payments, a healthy aircraft industry

is of‘ utmost importance to the United States. Because of this, the NASA COSAM

Program, industry efforts, and other activities underway are essential to

assure that if the cobalt supply is disrupted, viable substitutes and

alternatives are available to aircraft engine designers and producers.

METALLURGICAL CONSIDERATIONS

Cobalt is used in a variety of both cobalt-base and nickel-base

superalloys. The largest usage in terms of pounds consumed, however, is in

nickel-base alloys. Several nickel-base and cobalt-base superalloys are

listed in figure 6 which shows the range of cobalt content in these alloys.

Figure 6 also shows that cobalt is found in cast, wrought, and powder

metallurgy alloys. One of the largest consumers of cobalt is Waspaloy, a

wrought alloy used in aircraft engine turbine and compressor disks. These

components are heavy (thus consuming significant quantities of cobalt) but are

vital to today's high-temperature aircraft turbine engines.

14 . . In a recent publication by Tien, et al., a review of the existing

literature on the role of cobalt in superalloys was presented. In addition an

empirical model of the role of cobalt in superalloys was discussed at some

length. A summary of the various roles of cobalt on mechanical and physical

properties of superalloys observed by a number of investigators is presented

in figure 7. Based on a high-temperature torsion test, cobalt additions in

the advanced Nimonic series of alloys was reported to improve hot-workability.

P12-8
Figure 8, from the paper by Heslop (test temperature not stated),

illustrates that there is a linear dependence of the number of twists to

failure and the cobalt content of the Nimonic alloy.

16

Cobalt has been reported by Habraken and Coutsouradis to improve

oxidation/corrosion resistance of high temperature alloys. This conclusion

was based in part on the work on simple composition cobalt-chromium binary

alloys compared to similar composition iron or nickel-base alloys.

The primary role of cobalt in nickel-base superalloys of interest to

aircraft engine producers is its beneficial effect on high-temperature

strength properties. Figure 7 shows that cobalt's beneficial effect on

stress-rupture strength has been noted by Urbain^^ for Nimonic and Inconel

18

alloys. More recently similar results have been noted by Maurer

19 20

(Waspaloy), Nathal (MAR-M247), and Tien (Udimet 700) where cobalt was

systematically removed in incremental steps from these three complex

composition superalloys. The latter two research efforts are part of the NASA

COSAM Program and will oe presented during this Cobalt Session of the Workshop.

A summary of Urbain's ooservations is presented in figure 9. Stressrupture strength against temperature is shown for Nimonic 80A and Nimonic 90

(an alloy of similar composition to that of Nimonic 80A but with 17 percent

cobalt added). An improvement in stress-rupture strength of these wrought

alloys is noted as a result of the cobalt addition. A similar effect is noted

in figure 9 for tne cast alloys Inconel 713C (no cobalt) and Inconel 717C

(similar compositon, but with 8 percent cobalt added).

P12-9
Creep curves for reduced cobalt Waspaloy-type alloys, based on the work of

Maurer, are shown in figure 10. It should be noted that cobalt increases

rupture life, decreases minimum creep rate, and (negatively) reduces creep

ductility in Waspaloy at 730°C.

The effect of cobalt on microstructure of nickel-base superalloys is

summarized in figure 11. It should be noted that most of this research has

been reported by Heslop on simple composition nickel-chromium-cobalt

21

(60:20:20) ternary alloys. With the exception of the work by Lund , et al.,

very little work has been reporteu on complex composition superalloys. For

example, stacking faults were observed in the nickel-chromium-cobalt alloys,

while for a binary Ni-Cr alloy (80:20) stacking faults were not observed.

This lead to the conclusion that cobalt reduces the stacking fault energy in

nickel-base alloys. The effect of cobalt on aluminum and titanium solubility

in the nickel-chromium-cobalt system is shown in figure 12. The

microstructural influences of cobalt listed in figure 11 are interrelated to

the changes in mechanical and physical properties of the nickel-base

superalloys. However, the specifics of this interrelationship have not been

elucidated at this point in time.

SUPERALLOY FUNDAMENTAL RESEARCH

As was brought out in the preceding section of this paper, much of the

research aimed at determining the -role of cobalt in nickel-base superalloys

P12-10
has been conducted on simple binary and ternary alloys. As part of the: NASA

COSAM Program, the determination of the role of cobalt in today's more complex

nickel-base superalloys has been undertaken.

Five (four nickel-base and one cobalt-base) superalloys have been selected

for the COSAM investigation. The five alloys are listed in figure 13 along

with their typical applications in the aircraft engine industry, the forms in

which the alloys are useo, and remarks as to why they were selected for the

COSAM activity. Applications include turbine disks, turbine blades, and

combustors. A variety of product forms are represented by the applications of

the five alloys, as noted in figure 13. The selection of the five alloys was

based primarily upon the considerations given in this figure. Waspaloy was

selected because it represents the highest tonnage of cobalt in commercial

aircraft engines. Selection of Udimet-700 was based on the fact that this

alloy is used in the as-cast, as-wrought ingot, as-wrought powder, and as-HIP

powder-metallurgy fabricated conditions. The potential for determining the

impact of cobalt on both conventionally-cast as well as on single crystal

* turbine blades was the reason for selecting MAR-M247 . Rene' 150 was

chosen because it is one of the most advanced directionally-solidified turbine

blade alloys. The wrought, sheet alloy HA-188 was selected because it

represents one of the largest uses of a cobalt-base alloy in aircraft engines.

Trademarks

Waspaloy

Udimet

MAR-M

Rene

HA

United Technologies Corporation

Special Metals Corporation

Martin Marietta Corporation

General Electric Corporation

Caoot Corporation

P12-11
The primary purpose of the cobalt strategic element substitution research

is to determine the fundamental role of cobalt in a wide variety of

nickel-base superalloys and in a high-use cobalt-base superalloy. A secondary

purpose is to develop the methodology to explore the roles of other strategic

elements in similarly chosen alloys so as to have maximum impact on a wide

range of users.

Figure 14 shows the current participants in the COSAM activities on cobalt

strategic element substitution. These initial research efforts are planned

for a three-year period and consist of cooperative programs involving

universities, industry, and NASA Lewis Research Center. Nominal compositions

of the five alloys given in figure 14 indicate that cobalt content ranges from

10 percent in MAR-M247 to 39 percent in HA-188. In addition the y' phase

ranges from 20 percent in Waspaloy to 65 percent in Rene' 150. The first

phase in each research effort will involve substituting the less-strategicelement nickel for cobalt in incremental steps to a zero cobalt content. The

effects of this substitution on properties and phases present, such as y\

will make-up the major portion of the research effort in the first year of

each program element. Efforts in subsequent years will be directed at

identifying and optimizing alloying elements as substitutes for cobalt in the

five alloys so as to maintain the key properties of these alloys.

The cooperative nature (government-industry-academia) of the research

being conducted on Waspaloy and Udimet-70G is illustrated in figure 15. The

role of industry, as represented by Special Metals Corporation, is outlined.

Their primary role is to characterize and optimize fabrication and heat

P12-12
treating procedures for the reduced cobalt Waspaloy and Udimet-700 alloys.

The university role in this effort is also shown in figure 15. Columbia

University will be involved with mechanical property characterization,

structural stability, microstructural features, and theoretical formulations

to identify future alloy modifications if required for the second phase of the

project. Purdue University will be primarily responsible for microstructural

and microchemistry characterization of the reduced cobalt content alloys. To

round out the program, NASA Lewis Research Center will be involved in further

mechanical and physical metallurgy characterization of the alloys as shown in

figure 15. The output of this cooperative effort is expected to be a clearer

understanding of the role of cobalt in nickel-base superalloys to help point

the directions for future development of new/modified superalloys.

The research efforts on MAR-M247 and Rene' 150 parallel the previously

described efforts on Waspaloy and Udimet-700. It is anticipated that these

projects will lead to an understanding of the fundamental role of cobalt in a

variety of conventional and directional nickel-base superalloys. These

results should provide an improved technical base to guide the identification

of high potential modified superalloys in the planned COSAM Program.

CONCLUDING REMARKS

From the literature survey of the information on world cobalt producers

and the United States dependence on imports for over 90 percent of the cobalt

consumed in this country, it is clear that cobalt is truly a strategic metal.

Furthermore, the increase in tne use of cooalt for superalloys, which are used

P12-13
primarily for aircraft gas turbine engines, indicates the vital nature of

cobalt to this industry. Of particular concern is the impact that would

result to the aerospace industry if cobalt supplies were cut off for an

extended length of time.

A review of the literature indicates that the role of cobalt in

nickel-base superalloys is not clearly defined and the minimum amount required

has not been established. These factors give rise to the need for research to

determine the roles of cobalt in superalloys and to identify effective

alloying substitutes for cobalt. From these basic research studies, alloys

may be identified that iuinimize or eliminate the need for cobalt in

nickel-base superalloys, thus providing non-strategic element alternatives to

aircraft engine producers in the event of a cobalt crisis. The initial

efforts now underway in the NASA COSAM Program should help meet this national

need.

P12-14
REFERENCES

1. “Contingency Plans for Chromium Utilization", NMAB-335, National Materials

Advisory Board, National Academy of Sciences, Washington, D. C., 1978.

2. "Summary Report on Air Force Chromium Workshop", (AD-B015411L ) Air Force

Materials Lab, Wr ight-Patterson AFB, OH, Dec. 17, 1974.

3. Stepnens, J. R.; Barrett, C. A.; and Chen, W. Y. C.: "Corrosion/Oxidation

Behavior of Stainless Steels of Reduced Chromium Content", Reviews on

Coatings and Corrosion, Vol. 3, No. 4, 1979, pp. 211-252.

4. Dunning, J. S.: "Iron-Based Alloys Strengthened by Ternary Laves Phases",

Bureau of Mines, RI 8411, 1980.

5. "National Materials and Minerals Policy, Research and Development Act of

1980", Public Law No. 96-479.

6. Stephens, J. R.: "NASA's Activities in the Conservation of Strategic

Aerospace Materials", NASA TM 81617, Presented at 1980 Fall Meeting

of tne American Society for Metals, Cleveland, OH, October, 1980.

7. "GSA Seeks Cobalt Proposals", American Metal Market/Metalworking News,

April, 1981.

8. "Mineral Commodity Summaries", 1981, Bureau of Mines, Department of the

Interior, Washington, D. C., 1981.

9. "Cobalt, Mineral Commodity Profiles", Bureau of Mines, Department of the

Interior, Washington, D. C., October 1979.

10. "Weekly Petroleum Status Report", D0E/EIA-0208(81/18) , Energy Information

Administration, Department of Energy, May, 1981.

P12-15
11. Netschert, B. C.: "U.S. Dependence on Imported Non fuel Minerals: The

Threat of Mini-OPECs? " , Journal of Metals, March, 1981, pp. 31-38.

12. "Mineral Industry Surveys, Cobalt in 1978, 1979, and 1980", Bureau of

Mines, Department of Interior, Washington, D. C.

13. "Aerospace Review and Forecast 1979/80", Aerospace, Vol. 18, No. 1,

Winter, 1980, p. 4.

14. Tien, J. K.; Howson, T. E.; Chen, G. L.; and Xie, X. S.: "Cobalt

Availability and Superalloys", Journal of Metals, October 1980, pp.

12- 20 .

15. Heslop, J.: "Wrought Nickel-Chromium Heat-Resisting Alloys Containing

Cobalt", Cobalt, No. 24, September 1964, pp. 128-137.

16. Habraken, L.; and Coutsouracl i s , D.: "Tentative Synthesis on the Role of

Cobalt in High-Strength Alloys", Cobalt, No. 26, March 1965, pp.

10-24.

17. Urbain, M.: "Cobalt in Steels and Alloys", Cobalt, No. 12, September

1961, pp. 11-24.

18. Maurer, G. E.; Jackman, L. A.; and Domingue, J. A.: "Role of Cobalt in

Waspaloy", Superalloys 1980, J. K. Tien, ed., ASM, Metals Park, OH,

1980, pp. 43-52.

19. Nathal, M.: "The Role of Cobalt in a Nickel Base Superalloy", Presented at

the AIME Annual Meeting, February 22, 1981, Chicago, IL.

20. Tien, J. K.: "Role of Cobalt in Astroloy", Workshop on Conservation and

Substitution Technology for Critcial Materials", June 15, 1981,

Nashville, TN.

P12-16
21. Lund, C. H.; Woulds, M. J.; and Hockin, J.: "Cobalt and Sigma -

Participant, Spectator or Referee?", International Symposium on

Structural Stability in Superalloys, Seven Springs, PA, 1968, pp.

25-46.

P12-17
cn

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P12-18

AUSTRALIA
PRICE

TREND

COMPARISON

OF

COBALT

AND

OPEC

OIL

1972-1980

P12-19

COBALT
1980

UNITED

STATES

CONSUMPTION

OF

SELECTED

COMMODITIES

Q §

ii

P12-20
FIGURE

4

RECENT

TRENDS

IN

U.S.

AND

AEROSPACE

COBALT

USAGE

P12-21

1978

1979

1980
DISTRIBUTION

OF

1980

U.S.

COBALT

CONSUMPTION

(16

MILLION

POUNDS)

P12-22

OTHER
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P12-23

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P12-24

INCONEL

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P12-25

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P12-26

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P12-27

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FIGURE

12

COBALT

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TITANIUM-ALUMINUM

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P12-29

TITANIUM

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13

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P12-30
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14

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P12-31

CS-80-4369
FIGURE

15

COOPERATIVE

PROGRAM

TO

DETERMINE

FUNDAMENTAL

ROLE

OF

COBALT

IN

WASPALOY

AND

U-700

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P12-32

NASA

-

LERC

•

LOW

CYCLE

FATIGUE

•

THERMAL

FATIGUE

•

CYCLIC

OXIDATION

•

HOT

CORROSION
REDUCTION OF COBALT IN GAS TURBINE ENGINES

Dennis J. Evans

Pratt & Whitney Aircraft Group

P13

REDUCTION OF COBALT IN GAS TURBINE ENGINES

BY

DENNIS J. EVANS

PRATT AND WHITNEY AIRCRAFT GROUP

COMMERCIAL PRODUCTS DIVISION

EAST HARTFORD, CT

INTRODUCTION

When gas turbine engines were initially developed, many of the highly stressed

rotating or structural components were fabricated from high strength steels

or, where increased temperature capability was required, simple nickel-based

alloys containing chromium, and to a lesser extent aluminum or titanium and

carbon were used. The earliest of the "superalloys" to be used were Hastelloy

B (Ni-30%Mo-5%Fe-.12%Cl and Haynes Stellite Alloys 21 and 31, both cobalt

based casting alloys(l). Because of insufficient production capacity for

refining cobalt ores, an acute shortage of the metal occurred in the United

States at the time of the Korean conflict. This shortage spurred development

of nickel-based substitutes for the high-cobalt alloys and, when alloyed with

aluminum and titanium to form the strengthening compound gamma-prime (Ni 3

Al, Ti), led to such compositions as Waspaloy and Udimet 500.

Nonetheless, these alloys still contained 10 to 20 percent by weight of

cobalt, and consumed large quantities thereof, being widely used in disks,

blades, and to a lesser extent in vanes. Still preferred for use in vanes were

the cobalt-based alloys which possess superior thermal shock resistance. Thus,

with the advent of commercial jet travel, the consumption of cobalt in

superalloys for use in gas turbines increased to the point where in the late

1970's, it was the largest application for cobalt, exceeding even magnetic

materials and tool or die uses. The alloy Waspaloy achieved such widespread

utilization that it represents the highest tonnage use of cobalt in gas

turbi nes (2)

.

The political disturbance in Zaire, source of a majority all of the nation's

cobalt supply' served to sharply focus attention on the critical nature of

this element and its source of supply. Sharp cutbacks in material

availability, combined with increased demand, led to dramatic price

escalations beginning in 1977. Only recently has price stability been

restored, but at a level many times that in effect in 1975. Many efforts were

begun by industry, government and university researchers to find ways of

reducing use of cobalt in gas turbine engines. This paper will discuss three

approaches which have been successful in reducing consumption of cobalt-alloy

substitution, alloy modifi cati on and process development.

P13-1
Cobalt Applications in Gas Turbines

Three major sections of the gas turbine engine utilize the majority of cobalt.

These are the high pressure compressor, combustion chamber liners, and the

high pressure turbine. These areas are those portions of the engine where

highly stressed components operate at temperatures of about 850°F and above

- reaching to nearly 2200°F in nozzle guide vanes in the turbine section.

Such operating temperatures preclude the use of iron-based alloys and, until

just recently, all but the most creep resistant of titanium alloys.

The high pressure compressor utilizes static and rotating airfoils, usually

forged from nickel-based alloys with varying amounts of cobalt. Also found are wrought disk alloys, used to hold the blades. Similar wrought alloys are found

in the disks in the turbine section, and examples of these materials are grouped in Table I, together with their respective cobalt contents, under the

heading "disk alloys".

In the turbine airfoil area, where the highest temperatures are encountered,

the materials are usually cast, and consist of both cobalt and nickel-based

materials. Typical examples of this category of "cast airfoil alloys" are given in Table I, where it may be seen that Inconel 713 is unique in being

cobalt-free. Such was the case for Inconel 718 in the disk alloys. The

category of "other wrought alloys" has been reserved for those materials

typically processed into sheet, rod or bar forms for fabrication into

combustor liners, flame holders 'and other components exposed directly to the

combustion process, and thus requiring optimal corrosion, creep and thermal

shock resistance. The primary alloy used in such applications is Hastelloy X, so that consumption of cobalt in these components is not so large as in disks

and airfoils.

TABLE I

COBALT CONTENT OF TYPICAL TURBINE ALLOYS

Cast Airfoil Alloys

MARM 302

MARM 509

Rene 150

57 . 6% Cobalt

MARM 200 (Hf)

Inco 713

55.3

12.0

10.0

0

Disk Alloys

IN 100, MERL 76 18.5

17.0

13.5

9.0

0

Astroloy

Waspaloy

Rene 95

Inco 718

Other Wrought Alloys

L605 (Haynes 25) 53.4

HA 188

Hastelloy X 39.0

1.5

P13-2
Alloy Substitution

One of the earliest examples of alloy substitution to reduce cobalt usage is

the replacement of X-40 (Stellite 31) low pressure turbine blades with

nickel-base alloy blades. The Stellite 31 composition contains 56% (by weight)

of cobalt, plus 25.5% Cr, 10.5% Ni, 7.5% W, 0.5% C. In the early 1960's, the

cost of parts made from this air-melted cobalt alloy was less than for nickel

alloys which required the more costly vacuum melting and casting practice. As

can be seen in the cost trend information of Figure 1, the cobalt price

escalation of 1977 caused nickel alloy blade castings to now be less costly

than those of X-40.

Figure 1 Cobalt Price Increase Reversed the Relative Costs of Nickleand Cobalt-Based Cast Airfoils

A review of potential replacement alloys was made, with a total of five

identified. Considerable experience had been gained in each of the nickel

alloys considered - Waspaloy (13% Co) and Udimet 700 (18%) in wrought form,

and INIOO (15%), B1900 (10%) and Inco 713C (0% Co) in cast form. The cobalt

content of all these alloys was significantly less than that of X-40; however,

the complete absence of cobalt in Inco 713G, combined with the acceptable

level of mechanical properties (Table II) and in excess of 100 million

part-hours service experience led to the selection of Inco 713C as a substitute for X-40.

TABLE II

Comparison of Mechanical Properties Between X-40 and Inco 713C

Property Temp, °F X-40 IN713C

0.2% YS 1200 60 ksi 103 ksi

%EL 1200 8% 5%

Notched HCF (10^) 1200 24 ksi 31 ksi

Impact Strength 1200 7.0 ft-lb. 7.5 ft-lb.

0.1% Creep, 100 hr. 1400 20 ksi 60 ksi

Rupture, 100 hr. 1400 31.5 ksi 72 ksi

P13-3
A major step taken toward reduced cobalt^ consumption through alloy

substitution was the decision to incorporate Inconel 718 as a wrought disk

material to replace Waspaloy. Containing 13. 55^ cobalt, Waspaloy represents the

largest use of cobalt in commercial gas turbines. Millions of engine operating

hours of successful experience have been accumulated in a large number of

engine models. Because little experience in Inconel 718 wrought disks existed

at Pratt and Whitney, an extensive process and material evaluation program was undertaken.

Working closely with raw material suppliers and major forging vendors, two

components were selected for material and process qualification. One was a thin, compressor disk, while the other was the high pressure turbine disk for

the JT8D (Figure 2). Multiple heats of raw material were evaluated in each

configuration for each supplier to ensure ‘that the disks could be processed

repeatably, in a controlled manner.

Figure 2 Typical High Pressure Turbine Disk Used to Evaluate Inconel 718

Substitution

Extensive mechanical property specimen tests were conducted to generate design

allowables data in the area of tensile, creep, stress-rupture and fatigue. As seen in Figure 3, the tensile strength of Waspaloy is not significantly larger

than that of Inconel 718, and the latter material has a significantly improved

yield strength relative to Waspaloy. The acceptability of the tensile strength

has been demonstrated through overspeed testing of full scale disks.

Stress-rupture capability is of importance in components operating at elevated

temperatures such as the rims of turbine disks. Inasmuch as the critical areas operate at 1100-1150°F, the wide margin in rupture capability at 1250°F

(Figure 4) is not a serious drawback for most applications. In addition,

results of specimen and component fatigue tests have verified the suitability

of Inconel 718 for use.

P13-4
190

70 °F 1000°F 1200°F

Tensile

180 -

170-

160-

strength. 150 - ksi 140-

130-

120 -

110 -

100 L

UTS

.2YS

Waspaloy Waspaloy Waspaloy

IN 718 IN 718 IN 718

Figure 3 Tensile Strengths of Waspaloy and Inconel 718

300 hour

rupture

stress

Figure 4 Comparison of Typical Stress-Rupture Properties of Waspaloy

and Inconel 718

P13-5
Alloy Modif i cati on

As seen in the previous section, at temperatures above about IZOO^F, the

stress-rupture capability of Inconel 718 is sufficiently inferior to that of

Waspaloy that certain disks cannot be designed to make optimum use of Inconel

718. Better rupture capability is required. Studies both at Pratt and Whitney

Aircraft and Special Metals Corporation (3) have been directed at lower

cobalt modifications of Waspaloy. The objective of this research has been to

define an alloy composition with adequate design properties, especially

elevated temperature stress-rupture, while achieving reductions in cobalt

content

.

Several laboratory-sized heats of Waspaloy were prepared, with direct

substitution of nickel for cobalt. Variations of cobalt content were made from

the nominal 13.5% down to 8, 4, and 0% cobalt. Also varied was the amount of

carbon, since work of Betteridge, et. al. , on Nimonic 90 indicated an effect of cobalt on carbon solubility. After forging to small, flat

cylindrical shapes, which were then heat treated, an evaluation of tensile and

stress-rupture properties was made. While little effect on tensile properties

over the range of temperatures from 70-1200°F was noted, a strong dependence

of rupture life on cobalt content was noted at both 1200 and 1350°F (Figure

5).

Figure 5 Effects of Carbon Content on Stress-Rupture Properties of

Waspa loy

P13-6
Examination of the microstructure variations resulting from modifying the

nominal Waspaloy composition by replacing cobalt with nickel showed much the same trends as had been seen in the comparison of Nimonic 90 and Nimonic

80A(^). With successive reductions in cobalt content, additional carbide

precipitation was noted (Figure 6). This served to reduce grain size, which

may have contributed to rupture life decreases. More detailed microstructural

examinati on( 5) led to observation of changes in gamma-prime morphology, size

and distribution, as well as changes in stacking fault energy which

undoubtedly played strong roles in affecting high temperature properties.

Additional work in this area is in progress under a NASA-sponsored

program(6)

.

4% Co 0% Co

Figure 6 Reducing Cobalt Content in Waspaloy Tends to Increase Carbide

Preci pitation

P13-7
Net Shape Processing

The final method of reducing cobalt consumption has only recently begun to be

fully explored. Use of near-net shape processing approaches such as powder

metal disk fabrication by hot isostatic pressing and/or close-tolerance,

isothermal forging offers new avenues of reducing use of cobalt in gas turbine

components. As seen in the schematic disk cross section of Figure 7,

considerable amounts of raw material are utilized in producing a conventional

(non-isothermal ) closed-die forging, compared with the finish-machined part.

Figure 7 Cross Section of Typical Turbine Disk Forging Outline, with

Near Net Sonic Shape and Finished Part Outlines

By fabricating disks to the square-cut, rectilinear configuration in which the

ultrasonic inspection for quality control is conducted, significant amounts of raw material, and thus cobalt, may be conserved. In fact, as seen in Table

III, it may be possible to use an alloy with a higher percentage of cobalt,

and by appropriate net-shape processing, (^»^) achieve reductions of cobalt

of 10% or more.

P13-8
TABLE III

COBALT REDUCTION THROUGH NET SHAPE PROCESSING

A 1 loy Iput Wt. % Co Wt. Cl

Waspa loy 360 lbs. 13.5 48.6

Astro loy 250 17.0 42.5

Cobalt Savings = 6.1 lbs (12.5%)

Waspa loy 880 13.5 118.8

MERL 76 575 18.5 106.4

Cobalt Savings = 12.4 lbs (10.4%)

Future Opportunities

As can be seen from the examples cited in the foregoing section, substantial

progress has been made in reducing the use of cobalt in gas turbines. One

major area remaining, however, is the first stage turbine vanes, sometimes

referred to as nozzle guide vanes. Cast from cobalt-based alloys such as MarM509, these components can account for up to 40 pounds of cobalt in a typical engine. The cobalt alloys are used because of their high melting

points, resistance to thermal shock, and their good corrosion resistance.

Inasmuch as these properties are required mainly in the airfoil section of the

vane-that portion directly in the path of the hot combustion gases-it may be

possible to fabricate vanes using the critical materials only in selected

areas. Such a concept was discussed by Dr. J. R. Stephens (2) of the Lewis

Research Center of NASA in his presentation at the 1980 Fall Meeting of the

American Society for Metals. Dr. Stephens is the program manager for NASA's

COSAM (Conservation of Strategic Aerospace Materials) effort, which is

presently supporting cooperative research efforts in industry, university, and

government aimed at understanding the role of strategic elements and assessing

methods for minimizing their use. Tailored fabrication of airfoils may be one such approach to future reduction of cobalt, and other critical materials in

gas turbine engines.

P13-9
REFERENCES

1. Cobalt Monograph, ed. by Centre d' Infonnation du Cobalt, Brussels,

1960, pp. 255, 256.

2. Stephens, J. R.; NASA Technical Memorandum 81617, "NASA's

Activities in the Conservation of Strategic Aerospace Materials,"

prepared for the 1980 Fall Meeting of the American Society for

Metals, Cleveland, OH, October 28-30, 1980.

3. Maurer, G. E., L. A. Jackman, and J. A. Domingue, "Role of Cobalt

in Waspaloy," Superalloys 1980, J. K. Tien, et. al.. Editors, ASM,

Metals Park, OH, 1980, pp. 43-52.

4. The Nimonic Alloys, ed. by W. Betteridge and J. Heslop, Crane,

Russak & Co., New York, 1974, p. 78.

5. AFOSR Contract F49620-80-C-0048, "Reduction of Strategic Elements

in Turbine Disk Alloys," March 1980.

6. NASA Grant NAG3-57, "Role of Cobalt in Nickel-Base Superalloys," J. K. Tien, Columbia University, New York, New York.

7. NASA CR-135409, "Manufacture of Low Carbon Astroloy Turbine Disk

Shapes by Hot Isostatic Pressing," R. D. Eng and D. J. Evans, March

1978.

8. Contract F33615-77-C-5187, "Manufacturing Process for Production

of Near Net Shapes by Hot Isostatic Pressing of Superalloy Powder,"

Interim Report 14, February 1981.

P13-10
ROLE OF COBALT IN ASTROLOGY (SUPERALLOY)

John K. Tien and Robert N. Jarrett

Columbia University

P14

"Role of Cobalt in Astroloy (Superalloy)''

John K. Tien and Robert N. Jarrett

Columbia UniversityPresented at the Public Workshop on Conservation and Substitution Technology

for Critical Materials, 15-17 June 1981

1. Nickel-base superalloys are the single largest user of cobalt in

the United States.

2. The role of cobalt in superalloys is not well understood.

3. Data and experience base do not exist to warrant the amount of

cobalt currently in nickel-base superalloys in the United States.

4. Most of these alloys contain between 10-20% cobalt by weight.

5. Many Soviet and Chinese superalloys contain almost no cobalt.

6. In the past cobalt price was comparable to, and at times even less

than, the price for nickel. This might have been a major driving force in

the copious introduction of cobalt in the nickel-base superailoys in the

United States.

7. Columbia University (in conjunction with NASA-Lewis Research Center,

Purdue University and Special Metals Corporation) is in the first year of a

three-year program to understand the role of cobalt in nickel-base superailoys

in order to effect the prudent substitution of cobalt with such less critically

strategic elements as nickel.

8. The alloys chosen for research concentration are Udimet 700 (Astroloy)

with nominally 17-19% cobalt content and Waspaloy with 15% cobalt (the compositions of alloys found in Table I).

P14-1
2

9. Data Obtained so far strongly indicate cobalt has little effect

on the tensile strengths and ductility of either alloy; see, for example.

Figs. 1 and 2.

10. High temperature creep and stress rupture resistances are found

to be lowered as cobalt is removed from the alloys. However, this decrease

does not become significant until, for example, more than half the cobalt

is removed from the fine grained Udimet 700 studied; see, for example. Fig. 3

and Fig. 4.

11. Detailed and very fine spot size 'microchemistry determinations

in the scanning- transmission electron microscope show that cobalt rests

mostly in the continuous matrix part of the alloy and not in the strengthening precipitates; see, for example. Fig. 5 and Fig. 6. This can explain

why removing cobalt does not result in brute strength loss. The decrease

in creep and stress rupture resistances at very low cobalt levels can be

due to cobalt’s influence on the stacking fault energy of the matrix.

This decrease is not expected in every nickel-base superalloy, since many

nickel-base superalloys are not as dependent on their matrices as Udimet

700 and Waspaloy.

12. Results also show heat treatment can affect and under certain

conditions nullify cobalt's influence.

13. Tests in progress include determining the role of cobalt on

fatigue, thermal fatigue, oxidation and hot-corrosion, and other properties

relevant to the use of superalloys in jet engines and other critical applications .

14. A comprehensive report on the findings of this program is expected

by early 1982.

P14-2
TABLE

I

COMPOSITIONS

OF

THE

ALLOYS

STUDIED

(in

w/o)

LU rH LA rH CN rH

U_ rH 1-H r-H f-H

LP> cr CN NA 00

CQ CNI CN CN CN CN 0 0 0 0 0

CO rN CO CO CO 0 0 0 0 0 0

CN CT LA 0 CO

t-H rH 0 1 0

<C cr cr CT CT CT

LT\ CO LA CO . CO

1— NN NA NA hA NA

0 CT) rH 0 0 0 - - • - - 2: LA LA LA LA

a.

Q_

5: Q. CNJ Q.

2: 2: 2: O <

Q. Q- a. Q_ Q_ n_ V

O t-O O X _l

rH CNI (— H—

V V V > cQ O X CO OQ Q_

CH o

rH NA CO 00 0 0 - • - • •

C_3 CT 00 CN

V rH rH

rH CO CD CN . • • • • • z: CN Nn. NA CD LA

CO CO LA LA

0 CT LA CO CN NA

00 00 00 rA hA

00 CO 00 CD GO

h- rH 1-H 1^ rH f—

i

< 1 1 1 1 1 LU LA LA LA LA LA

x: Q Q Q Q Q

PlA-3
TENSILE

PROPERTIES

(25

C)

P14-4

COBALT

CONTENT

Cw/o)
FIGURE

2

-

Ambient

Temperature

Tensile

Ductility

of

Astroloy

Noiivwaod3a %

r

P14-5

1

COBALT

CONTENT

<w/o)
FIGURE

3

-

Creep

of

Cobalt

Modified

Astroloy

LlI

1—<

cr

QLlI

LU

cr

u

LDO

COBALT CONTENT <w/o)

P14-6

oo
FIGURE 4 - Stress Rupcure of Cobalt Modified Astroloy

STRESS RUPTURE (760*0

P14-7
FIGURE

5

-

STEPl/EDAX

^nalyais

of

ihe

Matrix

in

Modified

UDIMET

700

- «

. U)

. nj

- Ul

to

- n

- o


 <o

X Qi <t > Q. uj <r «-i r K uj z (/)-« I-

P14-8

UEIQHT

PERCENT

COBALT
FIGURE

6

-

STEM/EDAX

Analysis

of

the

Fine

Qanima

Prime

in

rVodified

UDIHET

700

X q: ^ > CLuj<iw »H r H Lj r -< I-

P14-9

UEIGHT

PERCENT

COBALT
PROGRESS IN REDUCTION OF COBALT IN nAR-,^247

Michael Nathal

Case Western Reserve University-NASA Lewis Research Center

P15
Progress in Reduction of Cobalt in Mar-M247

by

Michael Nathal

Case Western Reserve University

NASA Lewis Research Center

P15-1
INTRODUCTION

This work was funded by NASA Lewis Research Center.

The purpose of this investigation was to determine the influence of Co on

the microstructure and mechanical properties of Mar-M247, a nickel-base

superalloy. As one can see from the composition of Mar-M247, Figure 1,

superalloys are complex alloys, consisting of as many as a dozen elemental

adaitions to the nickel base. Throughout the years, an understanding of the

role of most of these elements has been reached. The role of Co, however, is

not well understood. Recent increases in price and limitations of supply

provide considerable incentive for understanding cobalt's functions, with the

goal of eliminating or reducing its role in superalloys.

The approach used in this study was to establish the properties of

equiaxed grained castings of the standard 10 percent Co version of Tlar-M247,

along with 0 and 5 percent Co versions. Nickel was substituted for Co to

produce the low 6o alloys. Grain size and heat-treatment were kept constant

prior to mechanical testing. Tensile tests were performed between 649°C and

982°C. Stress-rupture tests were performed between 760°C and 982°C.

P15-2
RESULTS

Microstructure

Figure 2 shows that Co has a significant effect on the microstructure of

Mar-M247. The figure on the left shows that the weight fraction of y '

>

determined by electrolytic phase extraction, decreases from 59 to 41 percent

as Co is removed. The figure on the right shows that the amount of carbides

is also a function of Co. The weight fraction of carbides increases from 2 to

3 percent as Co is removed. This increase in the amount of carbides promotes

the formation of a carbide film along the grain boundaries of the 0 and 5

percent Co versions.

The composition, as well as the amounts, of the phases are also a function

of Co level, as illustrated in Figure 3. The figure on the left shows the

composition of the y' phase as a function of Co level in the bulk alloy.

Important features of this figure is the increase in W and Ti contents in the

y' as Co is removed. Tungsten increased from 5 to 8 atomic percent, and Ti

from 2 to 3 atomic percent as Co was removed. The figure on the right shows

the composition of the y phase. Note the decreases in tne Cr and A1

concentrations in the y as Co is removed from Mar-M247

.

These changes in composition would serve to change the solid solution

hardening of the y and y' phases. The increases of W and Ti in the y' would

increase the strength of the y', mainly through increasing the APB energy.

The Cr and A1 decreases in the y would serve to weaken the matrix phase.

P15-3
Because Co is a weak solid solution hardener of Ni, the reduction in Co

content is not expected to alter the strength of either the y or y'* Thus, as

Co is removed from Mar-M247, the y' phase becomes stronger, but the y matrix

becomes weaker.

Tensile Properties

The summary of the tensile tests is presented in Figure 4. All three

alloys exhioit typical temperature dependence, with the yield strength

remaining fairly constant, followed by a rapid drop-off near 800°C. Note

that the 5 percent Co alloy showed a 12 KSI increase in yield strength over

the other two alloys, and the very small differences between the 10 percent Co

and the 0 percent Co versions of Mar-M247. Also of interest is the negative

shift in the peak strength temperature as Co is removed. This shift is

consistent with the increases in W and Ti concentration in the y' discussed

earlier.

Replotting the yielc strengtn as a function of Co level, Figure 5, again

shows the peak in strength at the 5 percent Co level. This peak is on the

order of 10 percent of the total yield strength, and is evident only at

temperatures below 760°C. Also evident from Figure 5 is tne very small

differences in strength between the 0 and 10 percent Co versions of Mar-M247.

In addition to the yield strength, the ultimate tensile strength exhibited

similar temperature and alloying dependence. Tensile ductility, as

illustrated in Figure 6, was not a strong function of Co level.

P15-4
Stress-Rupture Properties

Figure 7 is a summary of the stress-rupture tests, performed at 760, 871,

and 982°C.

In this plot of applied stress as a function of Larson-i^i 1 ler Parameter, a

reduction in life is evident at all temperatures. Rupture ductility was not a

strong function of Co, as illustrated in Figure 8.

0

Figure 9 shows the stress-rupture behavior at 871 C as a function of Co

level in Mar-M247. Rupture life is seen to decreased from 110 to 40 hours at

60 KSI, and from 265 to 90 hours at 52 KSI as Co is removed. These data are

consistent with the reduction in y' volume fraction discussed earlier, but

figure 10 shows that the creep behavior is not so simply related. This plot

shows that the minimum, or steatly-state creep rate, remains constant as Co

level is reduced from 10 to 5 percent, before increasing rapidly as further Co

is removed. This behavior is unusual, for alloys that exhibited a decrease in

rupture life, as shown in Figure 9, would be expected to exhibit a

corresponding increase in creep rate. This is seen to be true for the 0

percent Co alloy, where both the creep rate and rupture life are affected.

But this is not true for the 5 percent Co alloy, which exhibited a creep rate

equal to that of the 10 percent Co alloy. Another way to analyze these data

is to examine the creep curves produced, as in Figure 11. In this figure, it

is seen that both the 5 and 10 percent Co versions exhibit essentially the

same creep behavior until late in the test, when the 5 percent Co alloy enters

third stage creep earlier. The 0 percent Co alloy shows both the reduced

rupture life and increased creep rate compared to the 10 percent Co alloys.

P15-.5
This behavior can be rationalized in terms of the microstructural changes

in Mar-M247. One feature expected to be of importance in creep resistance is

the y' volume fraction, which was observed to decrease from 59 to 41 percent

as Co was removed, Figure 2. Such a reduction would be expected to increase

the creep rate and decrease the rupture lives of the alloys. This is seen to

be true for the 0 percent Co alloy, but not for the 5 percent Co alloy, and

other factors must be considered. Another feature considered to be of

importance in determining creep-resistence is grain boundary morphology. As

mentioned earlier, the 0 and 5 percent Co alloys exhibited a continuous grain

boundary film, while the 10 percent Co alloy exhibited discrete carbide

particles decorating the grain boundary. It is postulated here that the

presence of the grain boundary carbide film causes an increase in creep crack

initiation and propagation.

Concluding Remarks

In summary, the removal of Co did not significantly alter the tensile

properties of Mar-M247. The creep-rupture properties, however, were more

significantly affected. Removal of 5 percent Co reduced rupture life but did

not affect minimum creep rate. Removal of all 10 percent Co deleteriously

affected both creep rate and rupture life.

In light of these results, I believe that potential exists for use of low

Co versions of Mar-M247 in gas turbine engines, with modification. Two

modifications that seem to be promising are: 1. lowering the carbon content

to eliminate the carbide morphology problem, 2. using the alloys in single

crystal form, also to eliminate the grain boundary carbide morphology problem.

P15-6
MATERIAL: EQUIAXED GRAINED MAR-M247

Ni Co Cr Al Ti W Ta Hf Mo C

BAL. 10 8.5 5.51.0 10 3 1.5 0.8.15

NI SUBSTITUTED FOR Co TO PRODUCE

0% AND 5% Co VERSIONS

APPROACH: TENSILE TESTS BETWEEN 650°

AND980°C. STRESS-RUPTURE TESTS

BETWEEN 760° AND 980"C.

Fit'jre. I*

F(4ure J

AMOUNTS OF PRECIPITATES IN MAR-M247

AS A FUNCTION OF COBALT CONTENT

P15-7
ATOMIC

PERCENT

ELEMENT

COMPOSITION OF GAMMA AND GAMMA PRIME

AS A FUNCTION OF COBALT CONTENT IN THE BULK ALLOY

GAMMA PRIME GAMMA

Ficyjr^. Y

YIELD STRENGTH OF MAR-M247 WITH SEVERAL

COBALT LEVELS

AS A FUNCTION OF TEMPERATURE

YIELD STRENGTH

(KSI)

650 OC 750 850 950

TEMPERATURE cs- 81-729

P15-8
Y!ELu ilRLiYGi.i ;..i- M.SR-M247 FROM 649’ C (1200' F)

fO 982‘ C (1800' F)

S FUNCTION Of COBALT CONTENT

S- CS-81-739

f I qij ft (o .

TENSILE DUCTILITY OF MAR-M247 FROM 649° C TO 982° C

AS A FUNCTION OF COBALT CONTENT

20

649°C

DUCTILITY (%) 15

10 Q

5 _2_

1 L_

10

P15-9
LARSON MILLER PLOT FOR MAR-M2'47

WfTH 3 COBALT LEVELS

Fi Lt rd 7-

P/'c^ure S*

STRESS-RUPTURE DUCTILITY OF MAR-M247

AS A FUNCTION OF COBALT CONTENT

ELONGATION

(%)

REDUCTION IN AREA (%)

0 5 10 WEIGHT % COBALT

CS-81-2593

0 5 10 WEIGHT % COBALT

P15-10
STRESS-RUPTURE LIFE OF MAR-M247

AT 871* C (1600* F)

AS A FUNCTION OF COBALT CONTENT

RUPTURE

LIFE, HRS.

52 KSl

60 KSl

io.

STEADY STATE CREEP RATE OF MAR-M247 AT 871* C (1600* F)

AS A FUNCTION OF COBALT CONTENT

P15-11
MODEL CREEP CURVES FOR MAR-M247 WITH DIFFERENT

COBALT CONTENTS

|- I'g are / /

P15-12
REDUCTION OF STRATEGIC ELEMENTS IN

TURBINE DISK ALLOYS

FINAL REPORT

March 1, 1980-February 28, 1981

UNITED STATES AIR FORCE OFFICE OF SCIENTIFIC RESEARCH

BUILDING 140

BOLLING AIR FORCE BAS'E, D.C. 20332

PREPARED BY:

C. C. LAW, L. S. LIN, AND M. J. BLACKBURN

MATERIALS ENGINEERING & RESEARCH LABORATORY

PMTT&WHITLEY AIIRCRAFT GROUP

Government Products Division

p O Box 2691 West Beacfi Florida 33^02

UNITED

TECHNOLOGIES R P15-13
SUBSTITUTION FOR COBALT IN HIGH SPEED STEELSROLE OF PROCESSING INNOVATION

Walter T. Haswell Jr.

Colt Industries/Crucible Inc.

P16
CONSERVING COBALT

SUBSTITUTION FOR COBALT IN HIGH SPEED STEELS

ROLE OF PROCESSING INNOVATION

W. T. Haswell

Colt Industries/Crucible Inc.

P16-1
SUBSTITUTION FOR- COBALT IN HIGH SPEED STEELS - ROLE OF PROCESSING INNOVATION

The Crucible Particle Metallurgy process was developed by the

Crucible Material Research Laboratory and has been commercially applied

to the production of specialty steels at the Crucible Specialty Metals

Division in Syracuse, New York since 1970. This process enables us to

produce a superior quality product in standard analysis high speed

steels and also provides great opportunity for new alloy development

based on the virtual elimination of segregation made possible by the

CPM process

.

In mid-1978, the Katanganese invasion of Zaire seriously disrupted

fl the cobalt supply to the rest of the world. Cobalt became a scarce and

expensive metal. The United States imports about 977o of its cobalt

supply. This cobalt scarcity threatened to interrupt the production

of cobalt-containing high speed steels which are primarily used in

machining difficult- to-machine superalloys and titanium used by the

aircraft and aerospace industries.

At Crucible, we immediately started a research program to determine if through the unique segregation control of the CPM process we

could develop alloys offering equivalent metallurgical properties and

cutting tool performance to the cobalt-containing conventional super

high speed steels. This has been successfully accomplished and two

new cobalt-free super high speed steels, Rex 25 and Rex 20, are now

commercially available.

The Crucible Particle Metallurgy (CPM) Process

Before discussing the development of cobalt-free HSS, a brief

description of the CPM process is necessary. The principal elements of

the process are shown schematically in Figure 1. The basic process

P16-2
2

Particle Metallurgy Processing

j

U.-*-- SCREENING

1 r-i '

J _il iL iIil__J

HOT

ISOSTATIC

COMPACTION

Figure 1 - Particle Metallurgy Processing

P16-3
3

consists of induction melting and atomizing pre-alloyed particles,

screening and canning these particles, and finally hot isostatic

pressing to full density. The induction heat of the desired chemical

composition is melted and the molten stream is poured into an atomizing

chamber where high pressure gas jets disperse the molten stream into

spheroidal droplets which are rapidly quenched to ambient temperature.

The powder is removed from the atomizing chamber, dried and screened to

obtain the desired size fraction. It is then poured into cylindrical

steel cans which are evacuated and sealed. The cans are subsequently

heated to a specific temperature and hot i*sostatically compacted to

achieve a fully dense product. The compacts are processed to the desired

billet and bar sizes on conventional mill facilities.

Typical microstructures of conventional and CPM T15 high speed

steel are shown in comp,arison with conventionally produced product of

the same composition. Figure 2.

Figure 2

Conventional Product CPM Product

Primary carbides in microstructure of T15.

Photomicrographs taken at mid-radius location

on longitudinal sections of 2-in.-dia. bars.

Hardened condition .250X.

P16-4
4

The outstanding features observed are the homogeneity, particularly

in contrast to the segregation found in conventional ingot cast HSS,

and a fine uniform primary carbide distribution. The primary carbide

size distributions in CPM and conventionally produced T15 have been

measured. Virtually all carbides in CPM HSS have a diameter less than

about 3 whereas those in the conventional product cover the entire

size range to~^34 pm with a median size of 6 pm.

The principal advantages of PM HSS and the CPM process may be . summarized as follows:

(a) Superior grindability in the hardened condition.

(b) Improved toughness.

(c) Faster response to hardening heat treatment and better crosssectional uniformity of hardness.

(d) More uniform size change (less out-of-roundness) through

heat treatment.

(e) The ability to make more highly alloyed steel compositions

than is possible with conventional ingot cast product.

P16-5
5

CPM Rex 25 and CPM Rex 20 Alloy Development

Our alloy development program began with a review of the existing

high speed tool steel compositions to note the various combinations of

elements which are contained in currently available commercial grades.

The compositions were examined in two groups - the high tungsten

versions or the T series, and the high molybdenum versions or the M

series. Because of the wide compositional range represented by these

alloys, a common denominator was sought for comparison and analysis

purposes. The tungsten equivalency concept has been a popular means

of comparing HSS compositions and thus, in our review, we used the

relationship:

Tungsten equivalency, Weq = %W + 2 (7o Mo).

It was also realized that vanadium played an important part in

the alloy design concept, but for simplicity, our discussion will be

limited to the combination of W and Mo in the Weq concept. Table I shows the various grades which make up the T series. Beginning with

the grandfather of high speed steel, Tl, which contains 18W-4Cr-lV and

proceeding down the list to T15, it can be noted that 20% W represents

the upper limit of Weq which is contained in these alloys. Experience

has shown that this Weq level is the maximum which can be produced by

conventional ingot practice to obtain a sufficient yield to allow the

grade to be marketed profitably. Table II shows the various compositions which comprise the M series. Although there are many more grades

in the M series than in the T series, we again note that the maximum

tungsten equivalency contained in these grades is still only at the

20.5% level present in M33, M42 and M47.

P16-6
6

Table I. Compositions of Tungsten-Type High Speed Steels

Weight Percent

Type C W Mo Cr V Co Weq .

T1 0.75 18. 00 _ 4 .00 1.00 18

T2 0.80 18. 00 - 4 .00 2.00 - 18

T4 0.75 18. 00 - 4 .00 1.00 5.00 18

T5 0.80 18. 00 - 4 .00 2.00 8.00 18

T6 0.80 20. 00 - 4 .50 1.50 12.00 20

T7+ 0.75 14. 00 - 4 .00 2.00 - 14

T8 0.75 14. 00 - 4 .00 2.00 5.00 14

T9+ 1.20 18. 00 - 4 .00 4.00 - 18

T15 1.50 12. 00 - 4 .00 5.00 5.00 12

Table II. Compositions of Molybdenum-Type High Speed Steels

Weight Percent

Type C W Mo Cr V Co Weq .*

Ml 0.80 1.50 8.00 4.00 1.00 17.5

M2 0.85; 1.00 6.00 5.00 4.00 2.00 - 16

M3 Class 1 1.05 6.00 5.00 4.00 2.40 - 16

M3 Class 2 1.20 6.00 5.00 4.00 3.00 - 16

M4 1.30 5.50 4.50 4.00 4.00 - 14.5

M6 0.80 4.00 5.00 4.00 1.50 12.00 14

M7 1.00 1.75 8.75 4.00 2.00 - 19.25

MIO o 00 Ln 1— O O

- 8.00 4.00 2.00 - 16

M15+ 1.50 6.50 3.50 4.00 5.00 5.00 13.5

M30 0.80 2.00 8.00 4.00 1.25 5.00 18

M3 3 0.90 1.50 9.50 4.00 1.15 8.00 20.5

M34 0.90 2.00 8.00 4.00 2.00 8.00 18

M35+ 0.80 6.00 5.00 4.00 2.00 5.00 16

M36 0.80 6.00 5.00 4.00 2.00 8.00 16

M41 1.10 6.75 3.75 4.25 2.00 5.00 14.25

M42 1.10 1.50 9.50 3.75 1.15 8.00 20.5

M43 1.20 2.75 8.00 3.75 1.60 8.25 18.75

M44 1.15 5.25 6.25 4.25 2.25 12.00 17.75

M45+ 1.25 8.00 5.00 4.25 1.60 5.50 18

M46 1.25 2.00 8.25 4.00 3.20 8.25 18.5

M47 1.10 1.50 9.50 3.75 1.25 5.00 20.5

* Weq = Tungsten Equivalency, Weq - 7oW + 2(7oMo) .

*“ AISI types formerly listed.

P16-7
7

The most undesirable characteristic of conventionally produced

high alloy tool steels is the high degree of alloy and carbide

segregation which occurs during ingot solidification. This inherent

segregation is not only detrimental to the hot workability and fabrication of these alloys, but also results in a reduction in mechanical

properties and tool performance. An increase in the carbon and alloy

content above the levels contained in the currently available tool

steel grades results in an intensification of the segregation problem

and an intolerably low product yield after hot working. The Crucible

Particle Metallurgy or CPM process was developed to minimize alloy

segregation in standard high alloy tool steel grades, but as mentioned

above, a more significant feature of the CPM process is its use in the

production of more highly alloyed grades than can be made by conventional practices. The resultant CPM product is hot worked on conventional mill facilities and the carbide size and distribution are

controlled to develop optimized structures which offer significant

performance advantages.

Two prominent cutting tool grades used in machining the very

dif f icult- to-machine superalloys and titanium alloys used by the

aircraft and aerospace industries are T15 and M42 which contain 5 and

8/o cobalt respectively. Because of their importance in the machining

of materials used in our national defense effort, development of Cobaltfree alternatives for these grades appeared highly desirable from both a strategic materials and a marketing standpoint. Before our alloy

development program was initiated, the function of cobalt in high speed

steels was considered to enable us to better define the alloy design

concepts and property objectives which would have to be satisfied by a new alloy. Although the exact mechanism whereby cobalt enhances the

P16-8
8

properties of HSS has not been explained completely, certain general

property characteristics are normally attributed to the addition of

5 to 87o cobalt to HSS. In Roberts, Hamaker and Johnson's book, "Tool

Steels,"^ some of the beneficial effects of cobalt in HSS are listed

as follows:

1. Cobalt increases the solidus temperature thereby permitting

the use of high austenitizing temperatures to achieve greater

solutioning of alloying elements.

2. Cobalt enhances the secondary hardening reaction which results

in a 1 to 2 HRC hardness advantage in the fully heat treated

condition

.

3. Cobalt enhances the hot hardness and temper resistance

characteristics and thus allows a tool to retain a sharp

cutting edge at higher machining speeds that generate heat.

Thus, the objectives of the initial phase of our alloy development

program were to formulate a composition that would have room and

elevated- temperature hardness as well as temper resistance characteristics equal to those of T15. If these objectives were successfully

met, then impact and bend fracture strength, wear resistance and finally

tool performance properties would be compared. Although, at this point,

we realized that a cobalt-free HSS with these characteristics

could only be produced by our CPM (particle metallurgy process) because

of the anticipated necessity of increasing alloy content by a notable

amount, small laboratory ingots of various compositions were made to

expedite the hardening response and temper resistance studies.

Basically, our approach concentrated on the substitution of higher

tungsten equivalency levels for cobalt. One series of compositions

began with a cobalt-free T15 composition for a base alloy and

molybdenum was progressively added up to a tungsten equivalency level

P16-9
9

of about 347o. Samples of CPM Rex T15 were also included in our studies

to serve as a comparison base. The variation of heat treated hardness

and 1200F temper resistance with tungsten equivalency is shown in

Figure 3. The cobalt-free T15 analysis shows the predicted drop in

attainable hardness and temper resistance; however, both attainable

heat treated hardness and temper resistance showed a maximum value at

a tungsten equivalency of about 257o; hence, our designation of CPM Rex

25 for the optimized composition of this alloy development program.

Rex 25 heat treated hardness and temper resistance is comparable to

that of CPM Rex T15. Further alloy development work was conducted on Rex 25 using CPM heats until the carbon and alloy contents were optimized.

The final CPM Rex 25 composition is shown in Table III along with that

of CPM Rex T15.

Figure 3 - Effect of Weq on Attainable Hardness and

1200F Temper Resistance of CPM Rex 25.

Table III. Compositions of CPM Rex 25 and CPM Rex T15

Weight Percent

_C C^. _V_ . W Mo Co

CPM Rex 25 1.80 4.0 5.0 12.5 6.5 None

CPM Rex T15 1.55 4.0 5.0 12.25 None 5.0

P16-10
10

CPM Rex 25 v/as produced at the Specialty Metals Division in

Syracuse, NY, and processed to bar product to obtain material for

laboratory property comparison with CPM Rex T15. The results of these

comparisons are shown in the next four tables . Tables IV and V show

the results of temper resistance and hot hardness comparisons for

specimens which were heat treated to full hardness. These results show

that CPM Rex 25 and CPM Rex T15 are equivalent in both temper resistance

and hot hardness. Table VI shows the C-notch impact and bend fracture

strength values obtained on CPM Rex 25 in three heat treated conditions,

along with values obtained on CPM and conventional T15. Both the

C-notch impact and the bend fracture strength of CPM Rex 25 are slightly

lower than those of CPM T15, but are notably higher than those of

conventional T15 and were considered acceptable for this alloy.

Table IV. Temper Resistance of CPM Rex 25 and CPM Rex T15.

HARDNESS (HRC) AFTER INDICATED F.XPOSURES As Heat Treated

(2250F. 1025F/3X2 hr) Heat Treated + llOOF/2 hr nOOF/2+2+4 hr Heat Treated + 1200F/2 hr 1200F/2+2+4 hr CPM Rex 25 68 66.5 65.5 62.5 55 CPM Rex T15 67 65.5 64.5 61 54.5

Table V. Hot Hardness (HRC) of CPM Rex 25 and CPM Rex T15

Room Temp, Test Temperature Room Temp, Before Test lOOOF HOOF 1200F After Test

CPM Rex 25 67.8 57.8 54.0 48.4 64.3 CPM Rex T15 66.7 57.1 54.2 48.2 63.4

Table VI. C-Notch Impact and Bend Fracture Strengths of CPM Rex 25 and T15

Response to Heat Treatment and Properties -

* Aust

.

Temp, (F) HRC C-Notch Impact Strength (ft-lb) Bend Fracture StrenRth (ksi)

CPM Rex 25 2250 68 9.5 531 CPM Rex 25 2200 66.5 12.5 623 CPM Rex 25 2150 66 11 627 CPM Rex T15 2250 67 14 678

Conv, T13 2250 66 4 312

* A minute soak in salt bath and oil quenched. Triple tempered at 1025F

P16-11
Because of the favorable laboratory test results obtained on CPM

Rex 25, test material was supplied to about 30 toolmakers for use in

tool tests involving 14 different types of tools. The tests which have

been completed have shown CPM Rex 25 performance to be comparable or superior to that of conventional and CPM Rex T15. The highly favorable

field test results along v;ith a growing number of requests for bar

stock resulted in the commercialization of Rex 25 in December, 1979.

Results have been very encouraging and we are fully convinced that

another cobalt supply interruption would not present a problem for the

full replacement of T15 with CPM Rex 25.

Table VII. CPM Rex 25 Typical Field Test Results.

Tool Manufacturer

Application

:

Lab test drilling 1/2" holes in 1-1/4" thick 4340

HRC 31/33.

Results

:

1st Test 2nd Test 3rd Test

cp>; Rex M- 28 31 M42 17 23 _ CPM Rex 25 37 42 40

Tool Manufacturer Laboratory Testing End Mills Application: Profile milling 4340 at BHN 285 Results: Inches milled: 130" Wear land: T-15: 0.011" CPM Rex 25: 0,008"

Test bv an Aircraft Engine Manufacturer End Mills Application: Production milling a 286 high temperature alloy with

1/2", 4 flute end mill.

Results: CPM Rex 25 tool life performance 200/. better than M42,
12

When the CPM Rex 25 development effort progressed to the point

where laboratory properties obtained in the initial CPM experimental

compositions were favorable, a development program was begun on a cobalt-free alternative for M42 using the same principles and procedures as those used in the Rex 25 program. As producers of both

CPM and conventional M42, we knew that such a development would be

even more significant than one which provided a cobalt-free alternative

for T15 because the M42 market is notably greater than that for T15.

Consequently, successful development of this grade would result in a substantial reduction in cobalt usage in HSS.

A simplistic view of our design concept for this program would

be the addition of W to a 107oMo-17oV base. Again, a series of small

laboratory ingots were made to obtain material for a rapid assessment

of the heat treatment response and temper resistance characteristics

of various compositions. A cobalt-free version of M42 was included in

our experimental alloy series and shows the anticipated cobalt effect.

Throughout all of our studies, specimens of CPM Rex M42 were included

to serve as a comparative base for the experimental alloys. Figure 4

shows the effect of Weq on the hardening response and temper

resistance at 1200F.

H«rdn«*« -Rc CPM REX 20 DEVELOPMENT

A» H«at Traatad

2175*F 1025*x3

As Haat Traatad Abova + 1200*F 2 Hrs.

As Haat Traatad Abova

1200*F 2 + 2 hrs.

10 20 30

Turtgstan Equivalancy at 1.10X Vanadium

0N42

Q M42-no cobalt CPM REX 20 • Exp. 20 + 6% cobalt

Figure 4 - Effect of Weq on Attainable Hardness and

1200F Temper Resistance of CPM Rex 20.

P16-13
13

These results are similar to those obtained on Rex 25 except that the

maximum hardness was obtained at a slightly higher Weq level. A maximiam

hardness and temper resistance was obtained which was equal or superior

to that of CPM Rex M42. However, in our studies of the microstructures

of CPM versions of the Exp. 20 (I.IOV) compositions, we noted that the

grain size of Exp. 20 was coarser than we normally find in CPM product.

Therefore, CPM alloys with 2% V were produced and evaluated. The higher

vanadium version showed the desired grain size refinement. As shown in

Figure A, the heat treatment response and temper resistance of the 2% V

version was comparable or superior to that of M42.

To determine if it was possible to improve these properties, 67o cobalt was added to the Exp. 20 (1.107oV) base. It is interesting to

note that properties were not improved and thus the cobalt effect has

apparently been eliminated at this high Weq level. Additional compositions were evaluated in CPM form and the optimized CPM Rex 20

composition shown in Table VIII was selected.

Table VIII. Composition of CPM Rex 20 and CPM M42

Weight Percent

C Cr V W Mo Co

CPM Rex 20 1.30 3.75 2.00 6.25 10.50 None

CPM Rex M42 1.10 3.75 1.10 1.5 9.5 8.0

CPM Exp. 20 1.24 3.75 1.10 6.30 10.00 None

P16-14
14

A production heat of the CPM Rex 20 composition was melted and

rolled to bar to obtain material for property comparisons with CPM

and conventional M42, Table IX shows the 1100 and 1200F temper

resistance and Table X the hot hardness comparisons between CPM Rex

20 and M42 . These results show that our objective to attain

properties comparable to those of CPM Rex M42 has been fully satisfied.

Table IX. Temper Resistance of CPM Rex 20 and M42

Hardness (HRC) After Indicated Exposures

As Heat Treated

Hardness (HRC)

2175°/ HOOF 1200F

1025/3x2 hr. 2 hrs

.

2+2 hrs. 2 hr s . 2+2 hrs

,

CPM Rex 20 67.5 66 65.5 60. 0 57.0

CPM Rex M42 67.5 65.5 65.0 59. 0 55.5

Conv. M42 67.5 65.0 65.0 59. 0 55.0

Table X. Hot Hardness of CPM Rex 20 and M42

Hot Hardness (HRC) at Indicated Temperature

Room Temp

.

Test Temperature Room Temp

.

Before Test lOOOF llOOF 1200F After Test

CPM Rex 20 67.5 58.2 56.0 47.5 63.8

CPM Rex M42 67.0 58.4 56.0 48.0 63.0

Con"., . M42 66.5 58.5 56.1 48.0 62.2

P16-15
15

Table XI shows the Charpy C-notch impact and the bend fracture

strength results. CPM Rex 20 and CPM Rex M42 were equal in both

properties and both were notably superior to conventional M42 . In

^^*^ition, results are shown for the initial CPM Exp. 20 composition

which contains 1.14 V. The beneficial effects of the higher vanadium

content in the finalized CPM Rex 20 composition are obvious.

Table XI. C-Notch Impact and Bend Fracture Strengths of

CPM Rex 20, M42 and Exp. 20

Response to Heat Treatment and Properties

Aus t

.

Temp . op HRC

C-Notch Impact

Strength (ft-lb)

Bend Fracture

Strength (ksi)

CPM Rex 20 2175 67.5 12 581

CPM Rex M42 2175 67.5 12 581

Conv. M42 2175 67.5 5 372

CPM Exp . 20 2175 67.5 9 426

*1 minute soak in bath and oil quenched. Triple tempered at 1025F.

As a result of these favorable laboratory test results, CPM Rex

20 test material was provided to various toolmakers for performance

valuations. These field cutting tool tests were very successful in

that CPM Rex 20 provided equal or superior performance to M42.

P16-16
16

Table XII: CPM Rex 20 Typical Field Test Results

Tool Manufacturer Laboratory Testing

Application: End milling D2 tool steel, annealed

1/2" diameter, 4 flute end mills

Results

:

Length of Cut to .010"

Wear Land

CPM Rex 20 325 inches

Conventional M42

Production Test by

225

Form Tool User

inches

Application

:

Turning AISI S5 Tool

3" diameter circular

Steel

form tool

Results

:

Parts/Hour Parts/Tool Grind

CPM Rex 20 30 200

Conventional M43 30 100

Tool Manufacturer Laboratory Testing

Application: Drilling AISI 4150, 1/2" thick at Rc 33 hardness

1/4" drills

Results

:

Average Holes Drilled

CPM Rex 20 155

Conventional M42 84

The highly favorable field test results along with a growing number

of requests for bar stock resulted in the commercialization of CPM Rex

20 in March, 1981.

Results have been very encouraging and we feel that another

cobalt supply interruption or increased cost would not present a problem

in the replacement of MA2 and other cobalt-containing high speed steels

with cobalt-free Rex 20 or Rex 25.

P16-17
17

The Crucible Particle Metallurgy process has made it possible to

develop a new metallurgical principle on which a U.S. patent number has been

issued. This new alloying concept has led to the commercial development

of two cobalt-free super high speeds - CPM Rex 25 and Rex 20. The new

grades are equivalent or comparable to their cobalt-bearing counterparts

in the following properties:

1. Hardening response and attainable hardness.

2. Temper resistance.

3. Hot hardness.

4. Charpy C-notch impact strength.

5. Bend fracture.

6. Field cutting tool performance test.

P16-18
18

REFERENCES

1. W. T, Haswell, Jr., W. Stasko, and F. R. Dax, "Cobalt-Free

CPM High Speed Steels," The Metallurgical Society of AIME paper,

1980.

2. E. J. Dulis and T. A. Neumeyer, "Particle-Metallurgy of HighSpeed Tool Steel," pp 112-118 in Materials for Metal Cutting,

The Iron and Steel Institute, London, Publication 126, 1970.

3. A. Kasak and E. J. Dulis, "Powder -Metallurgy Tool Steels,"

Powder Metallurgy , 21 (2) (1978) pp 114-123.

4. A. Kasak, G. Steven and T. A. Neumeyer, "High-Speed Tool Steels

by Particle Metallurgy," SAE Paper 720182, 1972, pp 2-5.

5. John T. Berry, "High Performance High Hardness High Speed Steels,"

Climax Molybdemam Company, Greenwich, CT. 1970.

6. G. A. Roberts, J. C. Hamaker, Jr. and A. R. Johnson, " Tool Steels ,"

pp 710-713, American Society for Metals, Metals Park, Ohio, 1962.

The material contained in this paper is intended for general information only and should not be used in relation to any specific application

without independent study and determination of its applicability and

suitability for the intended application. Anyone making use of this

material or relying thereon assumes all risk and liability arising therefrom.

P16-19

W. T. Haswell

6/3/81
SUBSTITUTION FOR COBALT IN MAGNETIC ALLOVS

Thonas D. Schlabach

Bell Telephone Laboratories

P17
Substitution for Cobalt in Magnetic Alloys

by

T. D. Schlabach

Bell Telephone Laboratories

Murray Hill, New Jersey 0797^

The magnetics industry has been more successful perhaps

than any other in reducing its dependence on cobalt through

the use of substitutes that contain less or even no cobalt.

Thj-S is reflected in the fact that cobalt demand in this

area fell from nearly 25% of all cobalt demand in 1977 to

15% in 1980 according to U.S. Bu Mines data. Since most of

the magnetic use of cobalt is in Alnico V, the solution to our cobalt usage problem lay principally in finding substitutes for this alloy and, happily, several were available.

Beyond the permanent magnet alloys represented by the

Alnicos, however, cobalt finds wide use in other magnetically soft and semi-hard alloys for which substitutes were

also desired and were less obvious. Nonetheless, several

promising substitutes have been found in these categories

that contain no cobalt, are easier to process and that offer

substantial cost reductions with little or no sacrifice in

performance

.

_1_. Cobalt Substitution in Permanent Magnets

For many loudspeaker and magneto applications, particularly in consumer markets, product redesign to use cobaltfree hard ferrites (BaO*5pe^02) was possible and done.

Other possibilities for special use included cobalt rareearth magnets and Mn-Al-C, though the latter is not available commercially (1). Other promising permanent magnet

alloys include the low-cobalt Chromindur alloys, the Fe-Mn

and the Fe-Mo-Ni alloys shown in Table I that were developed

at Bell Laboratories for telecommunications use. Their use

in these applications provides a convenient means for discussing their properties and processing.

2 . Magnetic Alloy Substitution i n Telephone Receivers

Telephone receivers account for the largest single use of magnetic alloys and cobalt in the Bell System and the

current ring-armature receiver used is shown schematically

in Figure 1. Until recently, the permanent magnet in this

design was Remalloy ( Fe-20Mo- 1 2Co ) but developments started

in 1975 led to its replacement by the Fe-28Cr- 1 0 . 5Co alloy

known as Chromindur II. These Chromindur alloys, like

Alnicos, transform spinodally but, unlike Alnicos, are ductile. They can be processed to exhibit anisotropic magnetic

P17-1
2

properties by magnetic field aging, as is done with Alnicos,

or, more conveniently, by deformation aging. The latter

processing in the Fe-33Cr- 1 1 . 5Co alloy, known as Chromindur

III, leads to a maximum energy product ^^^rnax^» which

exceeds that of Alnico V at less than half its cobalt content .

In general , the Chromindurs offer magnetic properties

equivalent to the Alnicos at a third to a half their cobalt

content while combining good ductility and ease of processing with low cost.

Another important magnetic part of the receiver is the

ring armature currently made from 2V-Permendur having the

composition and properties shown in Table III. To perform

properly, the armature must have a high maximum permeability

^^max^> coercive force (H ) and good elastic properties.

The ability to substitute the Fe-6ri2Ni alloys for the 2VPermendur in this application is due to the development of a two-phase, low-temperature heat treatment that markedly

improves the soft magnetic properties of Fe-rich, Fe-Ni

alloys. These Fe-Ni alloys are also much easier to process

than 2V-Permendur and much cheaper (2).

3 • Reed Switches

Next to station apparatus, telecommunication's most

important use of cobalt-containing magnetic alloys is found

in reed switches that contain externally actuated contact

elements held in a sealed glass envelope. These remenant

reed sealed contacts currently use either Type 27 Remendur

or Nibcolloy as the reed element; both of which in Table II

are seen to have high cobalt contents. When the Fe-8Mn or Fe-7Mo-5Ni alloys are processed to exhibit an extremely

fine-grained, elongated and aligned, two-phase structure

they can offer very high remanence a desired range of

coercivity and high squareness (Bp/B^). These characteristics combined with good glass sealing and surface quality

make these alloys very attractive replacements for the standard semi-hard magnetic alloys (3,^). These substitution

possibilities are currently being explored.

PMT Memory

The permanent magnet twister (PMT) memory is a randomaccess, semi-permanent, nondestructive readout memory used

in electronic switching systems. It stores information in

the form of small Vicalloy permanent magnets whose properties are given in Table I. Like many other standard permanent magnet alloys, Vicalloy has a high cobalt content and

is a candidate for substitution. As can be seen in Table I,

both the Fe-l6Mn and Fe-20Mo-5Ni provide equal or higher

energy products in alloys that are easier to process, contain no cobalt and are much cheaper.

P17-2
3

These alloys are in the same families as those

described in Table II though processed in a different manner

to obtain permanent magnet properties (3,^). Beyond

memories, these new permanent magnet alloys are suitable for

speedometers, hysteresis motors, flag displays and antitheft tags for stores and libraries.

5 . Summary

The foregoing shows that the magnetics industry has

been able to respond quickly and effectively to the challenge posed by uncertainties in the availability and cost of

cobalt. This has been made possible by our deep understanding of magnetics and sophisticated thermomechanical processing of alloys, and further developments to eliminate or reduce our use of cobalt in magnetic materials may be anticipated. On the other hand, the unique properties conferred

on magnetic alloys by cobalt make its continued use assured

at a reasonable level.

P17-3
4

References

1.

Chin, G. Y. , "Impact of Recent Cobalt Supply

Situation on Magnetic Materials and Applications," IEEE

Trans Magn. MAG-15: 1685-91 (1 979 ).

2. Jin, S. , ^ "Ferritic Fe-Ni Alloys for Potential

Substitution of 2V-Permendur in Telephone Receivers,"

Presented at Int’l. Magnetics Conf . , Grenoble, France,

May 13, 1980.

3. Jin, S. , and Tiefel, T. H., "High Remanence Square-Loop

Fe-Ni and Fe-Mn Magnetic Alloys" IEEE Trans. Magn.

MAG-16: 1062-64 (1980).

4.

Jin, S. , and Tiefel, T. H. , "New Ductile Fe-Mo-Ni Magnet Alloys" to be pub. A ppl . Phy s . March 1981.

P17^4
PERf/AN'ENT MAGNET

ARMATURE

N

SUPPORT

Schematic illustration of the cross- sectional view of the U-type telephone

receiver.

Fig. 1.
TABLE I. New Permanent Magnet Alloys

Alloy (Gauss)

He

(Oersted)

max

MG'Oe

Fe-2 8Cr-10 . 5Co 9,800 400 1.6

Fe-33Cr-ll . 5Co* 12,000 780 5. 3-6 .

3

Fe-33Cr-7Co* 12,000 480 3.4-4 .5

Alnico V (24Co) 12,000 720 5.0

Fe-16Mn 10,000 270 1.7

Fe-20Mo-5Ni* 10,500 210 1.1

Vicalloy I (52Co)

*Anisotropic

9,000 300 1.0

TABLE II. New Semi-hard Magnetic Alloys

Alloy

Br

(Gauss)

He

(Oersted)

Br/Bs

Fe-8Mn 18,000 28 0.96

Fe-7Mo-5Ni

Type 27 Remendur

19,200 30 0.97

(49Co) 18,000 30 0.90

Nibcolloy (85Co) 15,000 20 0.95

TABLE III. Improved Soft Magnetic Alloys

Alloy ymax

Be

(Oersted)

B30

(Gauss)

P

( y^i-cm

Fe-6^12Ni 2,000 1.2 17 20

^^6,000 v2.

7

17 25

2V-Permendur (49C0) 4,500 2.0 20 26

P17-6
BORONIZING AS A SUBSTITUTE FOR COBALT BEARING ALLOYS

Stanley R. Scales

Hunhes Tool Co.

P18

BORONIZING AS A SUBSTITUTE FOR COBALT BEARING ALLOYS

BY

STANLEY R. SCALES

HUGHES TOOL COMPANY

I. INTRODUCTION:

ROTARY DRILLING CURRENTLY IS THE STANDARD PROCEDURE FOR MAKING HOLES IN THE EARTH'S

CRUST. ROTARY DRILLING IS USED FOR OIL, GAS, AND WATER WELLS, FOR BLAST HOLES IN

THE MINING INDUSTRY, AND FOR HOLES MADE FOR OTHER PURPOSES. CONTRIBUTING SIGNIFICANTLY TO THE SUCCESS OF ROTARY DRILLING HAS BEEN USE OF ROLLING CONICAL CUTTER

ROCK BITS. FIGURE 1 IS A BOTTOM VIEW OF A SOFT FORMATION ROCK BIT WITH THREE CONES

HAVING HARDFACED STEEL TEETH. THE CONES ARE ALSO CARBURIZED, HARDENED, AND TEMPERED. THE RING GAUGE IN THE FIGURE SHOWS HOW WELL THREE CONES FILL THE HOLE BEING

DRILLED AND PROVIDE THE MAXIMUM AMOUNT OF RELATIVELY LARGE TEETH AND BEARING CAPACITY.

FOR MANY YEARS, ANTI-FRICTION ROLLER BEARINGS, AS SHOWN IN FIGURE 2, HAD ADEQUATE

LIFE FOR ROCK BITS WITH STEEL TOOTH CONES. IN 1951, THE FIRST ROLLING CONE ROCK

BITS WITH CEMENTED TUNGSTEN CARBIDE TEETH WERE MARKETED. FOR SUCH BITS, THE TEETH

WERE CYLINDERS WITH HEMISPHERICAL CUTTING END, AS SHOWN IN FIGURE 3. THE CYLINDRICAL

PORTION OF EACH TUNGSTEN CARBIDE TOOTH IS PRESSED INTO A DRILLED AND REAMED HOLE

IN A NON-CARBURIZED LOCATION. ANTI-FRICTION ROLLER BEARINGS USUALLY ARE NOT ADEQUATE

FOR THE LONGER WEAR LIFE OF TUNGSTEN CARBIDE TEETH. JOURNAL BEARING, TUNGSTEN CARBIDE TOOTH BITS, AS SHOWN IN FIGURE 4, WERE INTRODUCED TO THE DRILLING INDUSTRY IN

1970-72, FOR THE FIRST TIME GIVING THE OILFIELDS A BEARING THAT LASTED AS LONG AS

THE CARBIDE TEETH. THE CONES FOR SUCH BITS WERE INLAID WITH A SOFT, ANTI-GALLING,

SILVER ALLOY IN THE GROOVES AND THE HOLES, AS SHOWN IN FIGURE 5, TO PROVIDE A SOLID

FILM LUBRICANT.^

P18-1
PAGE 2

THE BEARING PINS, AS SHOWN IN FIGURE 6, WERE HARDFACED WITH A COBALT BASE ALLOY,

OXYACETYLENE TORCH APPLIED.

BECAUSE THE PERFORMANCE OF JOURNAL BEARING TYPE, CARBIDE TOOTH. BITS WAS OUTSTANDING, INCREASED PRODUCTION WAS FORECAST AND HUGHES METALLURGISTS INITIATED WORK AIMED

AT FINDING SUBSTITUTES FOR THE SILVER ALLOY AND THE COBALT BASE ALLOY. COPPER ALLOY,

GTAW APPLIED, HAS REPLACED SILVER ALLOY IN MOST CONES OF THIS TYPE WITH SLIGHTLY

IMPROVED PERFORMANCE; AND BITS WITH BORONIZED BEARING PINS HAVE REPLACED SOME THAT

PREVIOUSLY USED COBALT BASE ALLOY HARDFACED BEARING PINS.^

II. BORONIZING OF BEARING PINS:

THE RESULTS OF LABORATORY BEARING MACHINE TESTS AND SURFACE TREATMENT TESTS REVEALED

THAT:

1. A CARBURIZED AND BORONIZED STEEL JOURNAL SURFACE ON A BEARING PIN, IF PROCESSED SO AS TO HAVE A SMOOTH SURFACE AND TO HAVE THE CARBURIZED CASE

CONVERTED TO TEMPERED MARTENSITE, PERFORMS SUPERIOR TO A SMOOTH HARD ALLOY

JOURNAL SURFACE.

2. A CARBURIZED STEEL SURFACE CAN BE COMPOUND BORONIZED IN A MANNER SIMILAR

TO COMPOUND CARBURIZING. FIGURE 7 SHOWS THE RESULTS OF TESTS USED TO

DETERMINE THE EFFECT OF TIME ON PACK BORONIZED CASE DEPTH AT THREE DIFFERENT TEMPERATURES.

FOLLOWING ARE THE PATENTED STEPS THAT CAN BE USED TO PROVIDE THE BEARING PIN OF A

ROCK BIT LEG (HEAD SECTION) WITH A CARBURIZED AND BORONIZED JOURNAL SURFACE, TYPICAL

MATERIAL BEING SAE EX30 STEEL:

P18-2
PAGE 3

1. COPPER PLATE THE HEAD SECTION EXCEPT FOR THE MACHINED BEARING SURFACES

OF THE BEARING PIN. THE COPPER PLATE PREVENTS CARBURIZATION OF NONBEARING SURFACES.

2. GAS CARBURIZE THE BEARING SURFACES OF THE HEAD SECTION BEARING PIN.

3. GRIND OR OTHERWISE POLISH THE CARBURIZED BEARING SURFACES AS NECESSARY TO

ASSURE THE DESIRED FINAL FINISH.

4. CUP PACK BORONIZE THE BEARING PIN FOR ABOUT 0. 005IN.(0. 1 3MM) BORONIZED CASE

DEPTH USING A MIXTURE OF POWDERED BORON CARBIDE AND POWDERED GRAPHITE,

WITH AN ENERGIZER ADDED. SEE FIGURES 8 AND 9.

5. UNPACK AND CLEAN THE HEAD SECTION, THEN QUENCH HARDEN AND TEMPER IT.

FIGURE 10 SHOWS A FINISHED BORONIZED HEAD SECTION READY FOR ASSEMBLY.

THE MICROSTRUCTURE OF THE JOURNAL SURFACE OF THE BEARING PIN OF A HEAD SECTION SO

PROCESSED IS SHOWN IN FIGURE 11. THE EXTREMELY HARD BORONIZED CASE IS SUPPORTED

BY A HARD, STRONG, TOUGH HEAT TREATED CARBURIZED CASE.

KNOOP MICROHARDNESS MEASUREMENTS (lOOG LOAD) ON THE BORIDE PHASES SHOW THE PRIMARY

F62 B PHASE TO BE 1550-1900 KHN AND THE TWO PHASE BORO-CARBIDE ZONE TO BE 1150-1450

KHN. IN CONTRAST, THE CARBURIZED CASE OF STEEL AT 60 HRC HAS A MICROHARDNESS OF

APPROXIMATELY 730 KHN.

P18-3
PAGE 4

III. PERFORMANCE:

A JOURNAL BEARING TEST HAS BEEN IN USE IN OUR LABORATORY FOR A NUMBER OF YEARS TO

MEASURE THE GALLING RESISTANCE OF MATERIAL PAIRS IN A LUBRICATED ENVIRONMENT. IN

THIS TEST, A BUSHING OF HARDENED AISI 4815 STEEL ROTATES ON A SHAFT MADE FROM

HARDENED SAE EX30 STEEL. THE DIAMETER AND CONTACT LENGTH OF THE SHAFT ARE SIMILAR

IN DIMENSIONS TO THE JOURNAL BEARING OF A 7-7/8 IN. (200MM) DIAMETER ROCK BIT. THE

INSIDE DIAMETER OF THE BUSHING IS CARBURIZED AND HARDENED (60 HRC) AND MAY HAVE SOFT

METAL INLAYS. THE SHAFT SURFACE IS THE TEST MATERIAL UNDER CONSIDERATION.

TESTS ARE CONDUCTED WITH LUBRICATED SURFACES USING A PRESSURIZED LUBRICATION SYSTEM.

LOADS AND RPM'S MAY BE ADJUSTED TO PRODUCE DIFFERENT PV (PRESSURE-VELOCITY) VALUES.

AT A CONSTANT RPM SETTING, LOADS ARE INCREASED AT HOURLY INTERVALS WHILE MONITORING

THE TORQUE NECESSARY TO ROTATE THE BUSHING. WHEN INITIAL SEIZURE (GALLING) STARTS,

THERE IS AN INCREASE IN THE TORQUE OF THE SYSTEM. TESTS ARE CONDUCTED UNTIL CRITICAL

SEIZURE (LOCKUP OF THE BUSHING AND SHAFT) OCCURS.

TABLE I SHOWS THE RESULTS OF PV TESTS CONDUCTED WITH SHAFTS (1) CARBURIZED AND

HARDENED (60 HRC), (2) CARBURIZED, BORONIZED, AND HARDENED, AND (3) HARDFACED WITH

A COBALT BASE ALLOY, STELLITE 190 (60 HRC). THESE RESULTS REPRESENT A LARGE NUMBER

OF TESTS CONDUCTED OVER A PERIOD OF SEVERAL YEARS. THE CARBURIZED STEEL-CARBURIZED

STEEL COMBINATION PRODUCES THE POOREST RESULTS. THE BORONIZED STEEL SHAFT EXHIBITS,

WITH BOTH CARBURIZED AND INLAY TYPE BUSHINGS, SLIGHTLY BETTER PV RESULTS THAN

STELLITE 190 RUN WITH A SILVER ALLOY INLAY BUSHING.

P18-4
PAGE 5

SUMMARY

:

THE USE OF BORONIZING FOR THE BEARING SURFACES OF JOURNAL BEARING TYPE ROCK BIT HEAD

SECTIONS HAS BEEN A SUCCESSFUL REPLACEMENT FOR COBALT BASE ALLOY APPLICATIONS.

FIGURE 12 SHOWS A BORONIZED HEAD SECTION AND A COPPER ALLOY INLAID, CARBIDE TOOTH

CONE REMOVED FROM A ROCK BIT AFTER 111 HOURS OF SERVICE.

P18-5
PAGE 6

REFERENCES:

1. J. R. WHANGER, U. S. PATENT 3,235,316, FEB. 15, 1966.

2. H. C. DILL AND S. R. SCALES, METALLURGICAL DEVELOPMENTS AID BITS, TOOL JOINTS,

OIL AND GAS JOURNAL (AUGUST 29, 1977).

3. S. R. SCALES, U. S. PATENT 3,922,038, NOV. 25, 1975.

4. S. R. SCALES, U. S. PATENT 4,012,238, MAR. 15, 1977.

P18-6
TABLE I

PRESSURE- VELOCITY (PV) TESTS FOR JOURNAL BEARING SHAFTS

PV AT INITIAL SEIZURE

SHAFT BUSHING PAIR PSI-FT/MIN MPA-M/S

CARBURIZED STEEL - CARBURIZED STEEL0.49 X 10^ 17.2

BORONIZED STEEL - CARBURIZED STEEL0.76 X 10® 26.6

BORONIZED STEEL - SILVER ALLOY INLAY0.84 X 10^ 29.4

STELLITE 190* - 0.58 X 10® 20.3

SILVER ALLOY INLAY

ALL CONTACT SURFACES ,GROUND TO A ROUGHNESS OF 32 X 10'^ IN

(8 X lO-^MM) RMS OR BETTER.

*NOMINAL COMPOSITION:

3.37%C, 26%Cr, 0.8%Si, 14%W, 6%MAX.Fe, l%MAX.Mo, 2%MAX.Ni

l%MAX.M,n, BALANCE Co.

P18-7
Figure 1

Botom View Of Three-Cone, Soft Formation Rock Bit With Ring Gauge

To Show How Well Three Cones Fill The Hole Being Drilled,

Gauge Diameter 200 mm (7-7/8 in.)

P18-8
Figure 2

Typical Roller Bearing For Ball Bearing Retained Steel Tooth Rock Bit Cone

P18-9
Figure 3

Tungsten Carbide Tooth, Hard Formation Bit Manufactured About 1960.

Note The Cemented Carbide Jet Nozzles Introduced In 1953.

Gauge Diameter 220 mm (8-3/4 in.)

P18-10
Figure 4

Tungsten Carbide Tooth, Medium Formation Bit Manufactured About 1972. One Bit

Third Has Been Sectioned To Show Journal Bearing For Ball Retained Cone. Gauge

Diameter 200mm (7-7/8 in.).

P18-il
Figure 5

Tungsten Carbide Tooth, Journal Bearing Type Rock Bit Cone With Soft, Antigalling,

Alloy Inlays In Bearing Surfaces.

P18-12
Figure 6

Bearing Pin Portion Of Head Section Used For Tungsten Carbide Tooth, Journal

Bearing Type Rock Bit Showing Cobalt Base Alloy Application On Bearing Surfaces.

P18-13
CASE

DEPTH

(in )

Figure 7

Boronized Case Depth As A Function Of Time And Temperature For Carburized SAE EX30

Steel

.

P18-14

CASE

DEPTH

(mm)
Figure 9

Head Section Of Figure 1 With Cup Packed, Ready For Boronizing.

P18-15
Figure 10

Head Section After Gas Carburizing, Pack Boronizing, Quenching, Tempering, And

Finish Machining.

P18-16
Figure 11

Microstructure Of The Boronized Case On Carburized SAE EX30 Steel.

500X

P18-17
Figure 12

Head Section And Cone From "Hughes" J22 Bit Which Drilled 264 Meters (1201 Feet)

Of Sand And Shale In 111 Hours At 20,250 Kg. (45,000 Pounds) Weight And 46 RPM.

Cone Inlays Were GTAW Applied Copper Alloy. Bearing Pin Was Pack Boronized.

Gauge Diameter 251mm (9-7/8 in.)

P18-18
TECHNOLOGICAL OPPORTUNITIES FOR A FIORE

STABLE SUPPLY OF TITANIUFl

Harris M. Burte

Air Force Flaterials Laboratory

V

P19

KEYNOTE ADDRESS - TITANIUM CONSERVATIONTECHNOLOGICAL OPPORTUNITIES FOR A MORE

STABLE SUPPLY OF TITANIUM

Harris M. Burte

Director Metals and Ceramics Division

AFWAL Materials Laboratory

The concern about disruptions or potential disruptions of the availability of titanium components in the United States is different than for

most of the other materials being considered at this workshop. There is no shortage of titanium reserves in the world; it is an abundant material.

Nor is there a shortage of titanium reserves in the United States. Although

the raw material for producing titanium metal is now primarily imported,

the amount of titanium tetrachloride produced from indigenous sources for

conversion to titanium dioxide pigment far exceeds the amount of titanium

tetrachloride which is converted into metal. Yet there is a concern about

increases in the cost of titanium metal and long lead times for the procurement of some products such as forgings. There is fear that imbalances

between supply and demand may continue to periodically reoccur. What is

the problem?

Fluctuations in the world market - between different users and uses

for titanium metal' and even potentially between metallurgical and pigment

use are a normal part of a competitive system. This workshop is not intended

to deal with these; rather it asks if there are problems unique or generic

to the titanium industry where technological advances may help insure a more

stable supply.

From its inception in the early 1950's, the growth of the titanium industry has had sharp ups and downs which have to now been tied to the fortunes

and mix of products in the aerospace industry. As the need for jet engines

grows or declines due to the overall needs of the military or commercial

aircraft market so has the outlet for a considerable fraction of titanium

production. As specific aircraft systems which use large amounts of titanium

(in addition to the engine) go into production (such as the Fl4 or F15) or are cancelled (such as the SST or Bl) periodic shortages or excesses of

capacity may occur. The expansion of titanium usage in the non-aerospace

market seems to have been relatively steady to the point where it has become

a significant fraction of production, but here too the possibilities for the

future introduce significant uncertainties. As the price of titanium com- ponents goes up will they maintain their hard won position in the chemical

Industry and in heat exchangers? Will major new uses evolve for energy

production or conversion, or in new marine systems? How good are projections

of future aerospace usage?

The inability to accurately forecast future demand may lead to recurring

periods of over capacity or shortage in the future, as in the past. Is there

* "Titanium" as used here also refers to titanium base alloys.

P19-1
one particular portion of the industry which is most sens! tive to this

situation? The "t i tan i urn" industry can be viewed as encompass i ng - four

major functions'':

mining and conversion to a chemical precursor such

as purified titanium tetrachloride.

reduction to titanium metal.

consolidation, e.g., arc melting to ingot.

production of sheet, plate, tubing, forgings, etc.

and conversion or assembly of these into finished

products

.

There seems to be general agreement that the first of these is not now, nor

is likely to become the bottleneck to titanium metal production. Titanium

tetrachloride incident to the production of pigment can be used for metal

with little extra effort. Which of the next three is the source of which

current shortages has been the subject of much controversy. At the precursor

workshop to this one 0) all were advocated with some fervor by representatives

of different segments of the industry. Minkler (2) suggests that from the

viewpoints of capital intensity and construction lead time, and since they

have utility for only one purpose, facilities for reducing tetrachloride to

metal are and will be the key step. A panel of the National Materials Advisory Board is currently making an assessment of "US titanium production

capability throughout the industrial complex from the source of raw materials

to the finished mill products" ( 3 ). A generally accepted identification of

current and potential future bottlenecks is obviously necessary and, although

the stated purpose of this workshop is to address technological options, I have included the following question for the discussion period to see if consensus can now be reached:

"What aspect of titanium production - reduction,

melting (consolidation), or mill shape fabrication (including forging) - is most likely to

encounter future inadequacies which can result

in significant shortages or price increases?"

From an overall viewpoint, will the "boom and bust" cycle continue, or will

the shakeouts of the past lead to an industry which is inherently conservative

and where new capacity tends to lag spurts in demand?

If history is any guide, and considering the uncertainties of use projections, periodic periods of shortage in titanium capacity is a reason- able future likelihood. There are two non-competitive approaches to try

to ameliorate such a situation. In the first the Government can decide if

it is in the national interest (e.g., for purposes of national defense) to

modify the current marketplace and/or how it operates to provide the stability

This categorization, although useful, does not fit all situations. Casting

for example, can combine consolidation with production of components.
or risk indemnification that will allow supply capacity to stay equal to or above needs. For example, the initial growth of the titanium industry

in the 1950's was due to Government encouraged financing, with loans and

stock-pile purchase contracts Under the current Defense Production Act,

the Government can direct production towards specific uses or can provide

incentives to creating additional capacity. Some have suggested direct

"Government purchases and/or guaranteed prices for certain metals such as cobalt and titanium'' . Additional options that are also currently

mentioned Include

replenishing the critical materials stockpile.

modification of tax laws to remove disincentives

to capital investment or to provide special

incentives in selected areas.

multi-year procurement for DoD programs.

increased fully funded advance procurement for

DoD programs.

The second approach considers the technological opportunities that

exist or might be developed to cope with the situation, and the identification of those is the primary purpose of this workshop. They include the

recognition and development of conservation and substitution options to

permit a temporarily inadequate titanium production capacity to be optimally

applied. From this point of view the options may be similar in nature (if

different in detail) to those presented for many of the other materials

being discussed here. They also, however, include opportunities which might

lead to less capital intensive production processes, to lower cost titanium

components and/or to expanded uses of titanium. The former might reduce

the risk of establishing new titanium capacity against anticipated growth of

need. The latter two by providing an expanded base for the industry reduce

the impact of the ups and downs in the less predictable portion of the market,

and also permit the most cost-effective uses to be more easily supplied during

times of shortage.

The expanded use of titanium in the areas it has already penetrated - aerospace, heat exchanger tubing and chemical process equipment - will perhaps be paced in the next decade more by reduction in cost than by the

development of higher performance alloys. In fact, if price escalation

increases at its recent rate, titanium may lose many of its hard won markets.

Titanium is also being considered or evaluated for a broad range of new applications These include steam turbines, ocean thermal energy

conversion, deep sea submers i bl es , geothermal energy and others. Successful penetration of such existing or potential markets may often require

improved materials as well as low cost processes.

Just as the nature of the concern about titanium is different from that

of many of the other materials being considered at the workshop, so will

3

P19-3
the types of technological opportunities which exist. Thus the primary

question I have posed for the discussion session is:

"What specific technological developments might

lead to significant reductions in the cost of

titanium components and/or expansion in its

non-aerospace use"?

I will now try to provide an outline of the different types of technological

development opportunities mentioned above and to illustrate with some examples of current or recent activity. I will indicate several papers in

other sessions which should interact with or contribute to ours, and I will

try to set a stage for our next two speakers. I will also offer a few

personal opinions about promising directions for R & D to stimulate your

response to the question posed above.

Subst i tut ion . Competition between materials or materials/design options has

been characteristic of the rapid expansion of materials technology in the

last several decades, and can be expected to continue (7). Titanium is no exception; its relative role will not necessarily remain static. If new competing materials replace it because they offer improved cost-effectiveness

they will do so irrespective of periodic shortages. Other competing materials

which involve some penality will, in principal, come into use only during

times of shortage or anticipated shortage. Selective use of such materials,

for applications which involve the least penalty, could help mitigate the

overall impact of shortages and the subsequent paper in this session on "Benefits of Titanium in Airframes and Compromises in Substitution" should

provide additional insight into such options. In my opinion it would not

be desirable to develop competing/substitute materials for titanium merely

to hedge against periodic shortages; the goals should be to surpass titanium

for specific uses and the efforts including expensive scale up and field

evaluation should be justifiable on this basis. The materials which result

would find specific uses on their own merits, and thus might be sufficiently

developed for other less effective use (involving some penalties) in the

event of titanium shortages. Some of the approaches which are now being or might be explored, which involve competition with present-day titanium technology,

are:

development of improved high temperature aluminum

alloys using rapidly solidified powder and/or

mechanical alloying approaches. (See the paper on "Potential for Rapid Solidification for Aluminum

Alloys" in Session 111 B).

development of high temperature advanced composites

such as graphite filaments in processable polyimide

matrices (see the papers in Session V).

improved design and use of advanced composites for

ambient temperature high load intensity applications.

A

P19-4
development of high fracture toughness high strength

steels for critical high load intensity applications.

It should of course be recognized that titanium technology need not sit

still, and the goals of the above activities should take this into con- sideration. Further, the competing approaches may involve their own jeopardy to materials shortages or price increases which, in a specific

case, may incur more risk than titanium.

Conservation . With titanium sponge approaching ten dollars a pound the

economic impetus for conservation becomes evident. Furtiier, if, as many

have suggested, reduction (spongemaking) capacity is most likely to be the

bottleneck in future shortages, conservation allows a limited supply to

satisfy more usage. However, as in the case of substitutes, I do not

believe that the development of conservation options can be justified solely

as a hedge against possible periodic shortages. Rather the cost of the R&D

and implementation of such options should be justified by the return on investment Irrespective of shortage considerations. The range of application of the technology so developed can then be extended if shortages

develop. Among the approaches, which have been or might be explored are;

near net shape processing - isothermal forging

(see the paper on "Potential for Near Net Shapes"

i n Sess ion 111 B)

.

progress is being made in developing casting processes

(and post casting treatments) and defining attractive

uses for the products; the long range potential of

developing new, compatible alloy/mold material combinations should be reexplored.

near net shape processing ' ' illurgy - recent

properties can be attained by a combination of attention

to cleanliness and control of the microstructure in

the final component. The paper on "Advantages in

Materials and Energy Utilization for Powder Metallurgy

Titanium" in this session will explore this subject

further (see also the paper on "Outlook for As-HIP'ed

Near Net Shapes" in Session 11 IB).

retirement for cause; the development of life prediction

methods to predict the growth of cracks under actual use conditions, of reliable nondestructive inspection methods

and of probabilistic decision models should permit the

safe use of components (such as engine discs) well beyond

the current practice of retiring an entire population at

the lifetime when the first one will have an unacceptable

ri sk of fail ure.

variable property components; the development of a variety

work in my own laboratory that rel iable

5

P19-5
of metallurgical synthesis methods will permit the

production of "one horse shay" components such as compressor discs with dual microstructures to yield

high fatigue strength at the hub and high yield

strength at the rim (9'.

repair technologies.

recycling; a closer evaluation of what impurity buildup

or other contamination is permissible for what uses would permit more progress, as would the development of

alloys/microstructures which are inherently more

tolerant to some contamination (which may be inevitable).

Although several of the approaches above do involve a significant "conservation"

aspect, they also contribute in other ways to cost reduction. In fact, these

other aspects may be their more important justification, i.e., the reduction

in machining cost for near net shape processes or the overall cost reduction

possible if casting a complex shape can provide acceptable properties.

Similarly, saving the overall manufacturing cost of most components is a more important aspect of "retirement for cause" than the value of the

sponge or other materials contained in them.

Cost Reduction . The prime impediment to significant expansion in the use and

range of applicability of titanium is cost, both the cost of sponge, and of

mill forms and final products. The aerospace experience (^9) js illustrative;

initial performance driven designs of many aircraft often specify much more

titanium than the final "cost affordable" product. Technological opportunities

for significant cost reduction can thus play the key role in expanding the

market for titanium and reducing the relative effect of inability to accurately

predict need in certain portions of the market. In some cases such cost

reductions may be necessary if existing applications are to be retained

against the competitive pressure of evolving substitutes.

Several of the approaches listed previously under

Conservation have already been identified as important more for their overall cost reduction

than for just the cost of the titanium saved. The

potential of powder metallurgy or casting remains

to be fully explored.

The extractive metallurgy (including reduction) of

titanium warrants considerably more long range/

innovative R&D attention than it has received in the

last two decades. Reducing the capital Intensity

and the energy cost of titanium are specific goals

to be kept in mind '^). Reduction of titanium tetrachloride directly to usable powder, and of mixed

chlorides or other precursors to alloy powders is a particularly intriguing approach. There have been

hints in the literature, but I feel it has not

received the attention it warrants. If attainable,

6

P19-6
it might offer the hope of low cost, high quality

input to the powder metallurgy fabrication approach

discussed elsewhere .

Are there approaches to producing large, high

performance titanium components which do not

require the large capital investment of heavy

presses?

Are there approaches for significantly reducing the

cost of machining titanium?

For aerospace applications significant effort has been

expended and progress is being made The

development of approaches such as superplastic forming

and diffusion bonding attempts to exploit certain

metallurgical characteristics of titanium and to

minimize the influence of others. Similar strategies

should be continued, and perhaps exploited to a wider

range of possible uses including the non-aerospace

market.

Titanium metallurgy is an actWe field in which significant progress is

being made. It offers many opportunities in addition to the few I have

mentioned. For example, I believe that rapid solidification approaches

(see Session 111 B),in particular, wi 1 1 open broad new vistas. I commend you all to a rededication of your energy, particularly to the

longer range, more Innovative possibilities which may emerge today or in

your future thinking.

I hope that the outline given above, and the possibilities suggested

will stimulate your input during the audience participation portion of

the work shop which will follow.

7

P19-7
REFERENCES

1. "Working Group on Critical Engineering Materials", Public Workshop

on Critical Materials Needs of the Aerospace Industry, Feb 9-10, I 98 I.

2. W. M. Minkler, "Titanium in I 98O", Journal of Metals, pp k]-kk,

Apr i 1 1981 .

3 . NMAB Newsletter, Jan I 98 I.

h. Chemical and Engineering News, pg 20, May 11, I 98 I.

5 . Testimony Before House Appropriations Committee, 13 Nov I 98O.

6. D. Eylon (ed) "Titanium and Energy", TMS-AIME Publication, in press.

7 . H. M. Burte, "The Nature of The Competition Between Emerging Materials"

Materials and Society pp 145-150 (1978).

8. D. Eylon & F. H. Froes "High Integrity Titanium Alloy Shapes by

Powder Metallurgy", presented at Symposium on Titanium and Its

Alloys for Surgical Implants.

9 . H. M. Burte and H. L. Gegel in "Process Modeling - Fundamentals and

Applications to Metals", pp 1-26, ASM Materials/Metalworking

Technology Series (I 980 ).

10. N. G. Tupper, J. K. Elbaum and H. M. Burte, "Opportunities for

Cost-Affordable Titanium Aerospace Structures" Journal of Metals

PP 7

- 13 , Sept 1978 .

11. F. H. Froes, D. Eylon, G. E. Eichelman, and H. M. Burte, "Developments

in Titanium Powder Metallurgy", Journal of Metals, pp 47-54,

Feb 1980 .

P19-8
ISOSTATIC PRESSING OF COnPLEX SHAPES

FRO!^ TITAN inn AND T I TAN I Un ALLOYS

Stanle^' Abkowitz

D'manet Tecbnolon'' Inc.

\ P20

ISOSTATIC PRESSING OF COMPLEX SHAPES

FROM TITANIUM AND TITANIUM ALLOYS

Stanley Abkowitz

Dynamet Technology Inc. Burlington, Massachusetts

The technology of cold isostatic pressing of powder metal into

complex shapes with enhanced mechanical properties is discussed. The

application of cold isostatic pressing, followed by vacuum sintering

to produce complex titanium alloy shapes with properties equal to

or exceeding those of castings in heavy to thin sections is demonstrated.

The significant economic and technical advantages of combining cold

isostatic pressing (CIP) with hot isostatic pressing (HIP) to produce

full density titanium alloy components are shown. The capabilities for

production of complex powder metal shapes such as impellers, and the

mechanical properties of these shapes achieved when manufactured from

titanium powder and alloy powder additions are demonstrated. Both cost

and energy saving advantages of the process are discussed.

P20-1
Introduction

I'old isostntic pressinq (CIP) of metal powders differs from

couvei'.tionai mecr.anical pressing in several ways that offer specific

advantages

.

(1) Since Hydrostatic (water-oil) pressure is applied uniformly on

all surfaces of the part to be consolidated, the density acliieved

IS of significantly greater uniformity.

(.2) with tnis uniform pressure tne size of tlie part to be compacted

IS limited only by the dimensions of the pressure ciiamber.'

(3) Since the mold material js c lat.tri

i

i

< e lyriacr' ‘i. pri'il. ard

d'- ! 1 ( lii.i;' ic 1 |.Ll'ican^. need bo aiUli'd to the piwder to minimize galling and die v;ear.

(4) With the elimination of lubricant no nre-'sintcr for lubricant

removal is necessary and no contaminati.on occurs.

(5) Witn tne use of elastomeric tooling, complex sliapes beyond the

' apab: 1 L :.j ec c-i ce nventional pressing are feasible.

In the past, isostatic pressing was employed to compact ceramic

r-.idterials at relatively low pressures and to form simple billets of

difficult to n\elt metals such as tungsten and beryllium. More recently

the technology has progressed to the manufacture of complex shapes

from a variety of metals.

The greatest economic advantages of isostatic pressing lie in Its application to metals that are relatively expensive and wiiich are ratner difficult to machine (since the process requires less

input material and reduces the extent of machining required). Certainly

one of the metals that fulfills these characteristics is titanium and

Its alloys.

As-Sintered Properties and Forging Preforms

Intensive developmental efforts of recent years ^2) have

shown that titanium and its alloys could be cold compacted and vacuum

sintered to achieve satisfactory density and useful mechanical properties.

This was accomplished using inexpensive elemental powders in combination

with master alloy powders as opposed to tlie more expensive pro-alloyed

powders, which in addition to their high cost are not readily cold

compactable

.

The tensile properties achieved with commercially pure titanium

and Ti-6Al-4V alloy are reviewed in Table I as follows:

P20-2
TABLE I. MECHANICAL PROPERTIES OF AS-SINTERED

AND SINTERED PLUS FORGED TITANIUM

C. P. Titanium (.12% Oxygen)

UTS

YSElong

.

R.A.

UTS

YSElong

.

R.A.

As-Sintered

(94% dense)

62.000 psi (430 MPa)

49.000 psi (340 MPa)

15%

23%

Forged Preform

( 100% dense)

66.000 psi (455 MPa)

53.000 psi (365 MPa)

*23%

30%

Ti-6A1-4V (.12% Oxygen)

As-Sintered

(94% Dense)

120.000 psi (830 MPa)

107.000 psi (740 MPa)

5.0%

8 . 0 %

Forged Preform

(100% dense)

133,500 psi (920 MPa)

122,000

psi (840 MPa)

11 . 0 % 24.7%

The as-sintered properties represent a 94% of theoretical density

product produced by isostatic pressing at 60,000 psi (415 MPa) followed

by vacuum sintering at 2250°F (1232°C) for 3 hours. An improvement in

properties is achieved by forging a 94% dense preform. The as-sintered

tensile properties compare well with those properties achieved by

conventional casting and offer an economic process for li)ce applications.

Where higher density and higher mechanical properties are required,

forging the pre-shaped P/M preform may be employed. Other metal worl^ing

techniques which have been proven to fully densify the as-sintered

preform are rolling and extrusions. Typical forging preforms and a finished forged part are shown in Fig. 1 and Fig. 2.

By designing the elastomeric tooling in conjunction with a mandrel,

cored shapes can be produced with configurations pressed and sintered

close to size. The large titanium valve ball of Fig. 3 and the titanium

alloy liner of Fig. 4 were produced by isostatic pressing around a mandrel which was removea from the cold pressed compact prior to vacuum

sintering. The rotor cannister shown in Fig. 5 demonstrates the

producibility of very thin walls (.035 inch/. 90 mm). These parts are subsequently finish machined to final dimensions.

In the above applications the components possess the desirable

mechanical properties and corrosion resistance in the as-sintered

condition. It is interesting to point out that in addition to the

considerable material and machining time saved by the preformed shape,

further savings are offered in the finish machining operations which

still must be performed. This is due to the improved machinability of

the P/M material as indicated in Fig. 6. At 100 hole "tool life" the

cutting speed for annealed Ti-6Al-4V improves from 25 feet per minute

for the wrought product to 60 feet per minute for the P/M material .

P20-3
FIG 2 TITANIUM COMFRIOSSOR BLADE FORGING P/M T1-6A1-4V PREFORM

AND FINISHED FULL DENSITY FORGING

P20-4
FIG. 4 T1-GA1-4V LTMER 8" (203 mm) DIA. AS PRESSED AND SINTEPXD

P20-5
FIG. b Ti--GA1-4V GOTOR C/ANNIS^IjR AS PRESSED A:jD SINTEREL)

_.^0

I hOLLKO. solut.on I

u 2U '<b uU lob rurfiNCi L) II i.i/.viii.iiii

FU;. t,

lAClII’IAlilEITV STUDY ON P/il •m’ANUlM iiAl 4V ;DE3Y

P20-6
Subsequent Hot Pressing

For certain applications where full density is required, the hot

working operations of forging, extrusion, or rolling can be performed

on P/M preforms, but where deformation is not required, hot isostatic

pressing (HIP) can be used to produce fully dense shapes . This is

demonstrated by the microstructures of Figs. 7A and 7B which show the

change in porosity of Ti-GA1-4V.

The clu;mical analysis and mechanical properties achieved alter llIP are shown in Table II below and compare favorably with conventional 100%

^ dense wrought product.

TABLE II. MECHANICAL PROPERTIES t. CHEMICAL ANALYSIS

( OF A Ti-GA1-4V SINTERED AND HIP PREFORM

TENSIU; TEST AT ROOM TEMP .

:

CHEMICAL ANALYSIS:

Tensile P.S.I. 132,797 (916 MPa) Iron .085

Yield .2% offs. 120,024 (828 MPa) Vanadium 4.07

Elongation in 4D 13.0% Aluminum 5.97

Reduction Area 26.0% Carbon .035

Hydrogen - Core .0018

Nitrogen .0109

Oxygen .1978

Tnis combination of cold isopressing. sintering and hot isostatic

pressing is a highly economical technique for achieving maximum density,

particularly for titanium alloys. The procedure is often referred to as CHIP. The prior cold isostatic pressing and sintering permits use of

elemental master alloy powder which is significantly less expensive than

pre-alloyed powder, as well as the use of elastomeric tooling during cold

compaction to shape rather than expensive and expendable hard tooling

used for canning pre-alloyed powders. In this manner, the sintered

preform can be HIP as inexpensively as commercial titanium castings are currently densified. Rotating beam fatigue endurance limits at 10^ of

52,000 psi {360 MPa) have been obtained with Ti-6A1-4V processed by this

technique

.

Complex Shapes

The technology of both CIP and CHIP processing permits significant

advantages in the manufacture of complex shapes such as impellers.

Figs. 8 and 9‘ show the detail that can be acnieved in the complex

configuration of a radial impeller. Cost savings in the manufacture of

such shapes can be as high as 70% compared to conventional machining from

billet or rough forging, even where subsequent machining might be required.

Fig. 9 indicates the potential (with recent developments in tooling) for

production of multi-holed structures and perforations in both thin and

thick sections. Tl»c titanium alloy disc shown has over 1000 hoU's and is 3/L6 inch (4.76 mm) ttiick. Heavier sections with square, trianciular, or any shaped hole could readily have been incorporated into the tooling.

P20-7
FIG. 7A

HICROSTRUCTURE Ti-oAl-4V

AS PRUSSED AND SINTERED

FIG. 7B

iicrostructuft; ti-gai-jv

COLD ISOSTATIC PRESSED AND SINTERED FOLLOWED

3Y HIP

DENSITY

\V-w Ai

...

I 500X

T1-0/M-4V CHIP ».

'•/’JK r-. Ti-6A1-4V CHIP .5 qQ4.% nrMCT’pv 'i'f-* .i * 4. 1’

I wS rvV'- '

P20-8
FIG. 10

Ti-6A1-4V NOZZLE PREFORM AS ISOPRESSED AND AFTER VACUUM SINTERING

P20-9
Cost and Energy Savings

Cost saving applications have emerged for isostatic pressing from

industries as diverse as instant photography and missile manufacture.

Fig. 10 illustrates a Polaroid titanium nozzle section where very

expensive machining operations were previously employed. The use of

isostatic prcssiiKj witli a detailed liard mandrel produced a net sliapo

for tile internal configuration and thus eliminated the major costs of

machining these components.

In the missile wing configuration shown in Fig. 11 which is currently under development, similar cost savings are forseen with the

advantageous application of elastomeric tooling in combination with

mandrel technology.

One component currently in production at the rate of 300 parts

f>er month is a titanium alloy Dome Housing for the Sidewinder Missile

;See Fig. 12). This item had boon machined from a billot blank weighing

over 5 lbs. The final shape is machined to close tolerances and a very

thin wall of .025 inch (.635 mm). Application of isostatic pressing

and vacuum sintering to this component resulted in a preform blanic weighing 1-1/4 lbs. and significantly reduced t)ie amount of finish

machining necessary. In addition, the final machining required was on a free machining material, thus enabling the close tolerances to be more reaaily achieved.

In many cases the cost of the P/M preform might well exceed the cost of ttie billet stocl:, with significant eventual cost savings still

achieved in reduced finish machining. In the specific case of the housing

however, tne cost of the preform is only 60% the cost of billet, with

additional still better savings in finish machining costs.

Both cost savings and energy savings can be obtained through the

process described here. Previous manufacture of the Dome Housing from

"mill product" required the use of 5 lbs. of material which had to be

extracted from ore, reduced and refined, melted and forged.

Since the P/M process starts with powder produced in the manner of

sponge, and since the preform blan)c uses only one forth as much material,

the energy requirements for extraction, reduction and refining are reduced to 25% of the energy needed for conventional processing. In contrast to the significant energy required to vacuum double arc melt

and forge to billet 5 lbs. of material is the much reduced energy to press and vacuum sinter only 1-1/4 lbs. to preform shape. Finally, even

finish machining to final dimensions of both the wrought and P/M preform

results in still additional energy savings due to the free machining

aspect of the material.

Other Materials

The technology above has been applied to materials other than

titanium (4) (5)Stainless steels, tool steels, aluminum alloys,

copper, nicltel, Hastelloy C and the refractory metals among others

have all shown technical promise and economic advantage witli these

processes. Increasing comn\ercial acceptance and significant utilization

of this tcclmology is anticipated in the immediate future.

P20-10
FIG. 11

T1-6A1-4V MISSILE WING ISOSTATICALLY PRESSED AND VACUUM SINTERED

FIG. 12

T1-6A1-4V ALLOY DOME HOUSING PRllFORM LEFT FINISHED MACHINED COMPONENT RIGHT

P20-11
References

li) S. Abkowitz, J. M. Sierjiej, and R. D. Regan "Titanium P/M Preforms,

Parts and Composites". Proceedings of 1970 international Powder

Metallurgy Conference, Vol. 4, Pages 501-511.

(2) S. Abkowitz "Cost Savings in the Application of P/tl Titanium and

P/M Aluminum Alloys". 1974 National Powder Metallurgy Conference

Proceedings, Vol., 30, Pages 85-101.

(3) Metcut Report to Dynamct Technology "Drilling Pressed and Sintered

Titanium 6A1-4V". Metcut Report No. 1023-12894-1, February 1969.

(4) "Developments in Titanium and Aluminum P/M Technology", Precision

Metal Magazine, September 1974, Pages 44-48.

(5) "Cold Pressing Opens the Door to Complex Shapes", Iron Ago Magazine,

.'lay 9, 1977, Pages 39-41.

(6) S. Abkowitz, "Isostatic Pressing of Complex Shapes from Titanium

and ether Materials" ICt ; National SAilFE TECHNICAL CONFFPirNCE , Kiamesha

r.ake. New York, 1978.

P20-12
BENEFITS OF TITANIUN IN AIRFRAMES AND

' COMPROMISES IN SUBSTITUTION

Donald G. Goehler and A. Lew Wingert

Boeing Commercial Airolone Co.

P21
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senefits of titanium in airframes and

COMPROMISES IN SUBSTITUTION

Donald D. Goehler and A. Lew Wingert

Manager, 757 Materials Technology

Boeing Commercial Airplane Company

ABSTRACT

The Boeing Company has used titanium from the time it became commercially available to enhance the performance of our products. Each new product

has used more titanium than earlier products, and it is expected that this

trend will continue. Supersonic cruise vehicles will require that substantial amounts of their airframes be fabricated from titanium to optimize

their efficiency. Substitution of other metal alloys for titanium in

existing and some future application extracts a penalty in performance.

INTRODUCTION

Titanium has been the subject of research, development, and application

at The Boeing Company since the late 1940s. Its properties of high

strength, low density, and corrosion resistance make it very des.rable as an

airframe material. Initial efforts by our manufacturing department to make

airplane parts, however, were marked by high cost and limited success. The

titanium industry was encouraged to invest private capital in equipment and

techniques that would give us raw material forms that we could easily shape

into parts for air vehicles. We invested our own research money and

participated in government research to form, weld, fasten, cast, and machine

titanium. Slowly, commercially pure titanium, then Ti-6A1-4V, found

increasing applications, as shown in Figure 1. Starting with the WS-llOA

P21-1
(eventually called the B-70), then the supersonic transport, Boeing encouraged the titanium industry to build additional production capacity using

their own capital. Figure 2 shows the results of program cancellations

on titanium production and shows why that industry has difficulty funding

expansions.

These airplane programs, even though cancelled, made us more familiar

with the characteristics of titanium. We learned that strength at elevated

temperature was unmatched by any other metal of the same or lower density.

Its corrosion resistance, comparable to stainless steel, meant that parts

could be used in corrosive environments, such as landing gears, without

special plating processes. We also remained firmly convinced that titanium

was right for any supersonic cruise air vehicle.

Although we still use commercially pure titanium for ducts and tubes,

our interest and major uses turned to the alloys of titanium, particularly

Ti-GA'-TV, which has served as a workhorse for the aerospace industry since

the late 1950s. Its moderately high strength, 120,000 psi ultimate tensile

made it competitive with the high-strength aluminum alloys in a

strength/density comparison. In the middle 1960s, we applied Ti-6Al-6V-2Sn

at 160,000 psi ultimate tensile strength, which allowed titanium to compete

directly with ultra-high-strength steel alloys in landing-gear applications

Our newest airplane, the 757, is applying Ti -10V-2Fe-3Al at 180,000 psi

ultimate tensile strength, which exceeds the strength/density of ultrahigh-strength steels (Figure 3). Substituting this alloy for Ti-6A1-4V in

our main landing gear beam saved approximately 125 pounds per airplane.

To illustrate the expanded use of titanium at Boeing, the amount of

titanium in the airframe of the 727 is less than 1 percent (1300 pounds),

whereas there is 4-1/2 percent (11,000 pounds) titanium in a 747 airplane.

When you include the engines, we buy about 72,000 pounds of titanium raw

material for every 747 we build. The 757 will have 6-1/2% titanium in

P21-2
its airframe when it rolls out in December, 198L

From the above, it can be seen that titanium technology is a vital

tool for Boeing when building competitive products in a world aerospace

market. Our own up and down requirements, the supply complexities, and

our need to reliably build products on schedule cause us to examine the

process of decision making when selecting or substituting materials.

THE MATERIAL SELECTION PROCESS

Thousands of highly trained people, with the best tools in the world,

highly experienced in the many complex disciplines required to thrive in a

high-technology, worldwide enterprise, spend about 4 years to develop, sell,

design, procure, build, test, and deliver a new airplane. In production,

some components, like the 747 main landing gear beam (Ti-6A1-4V), must be

ordered 4 years before we install them in the airplane. Therefore, the

decisions about what materials we use to build a Boeing airplane must be

made reflecting all of the business opinion and supply questions, as well as

airplane design efficiency.

When it is decided what materials we will use, it requires fast

footwork and shrewd negotiations all over the world by out procurement

specialists to ensure that those materials, meeting our specifications,

arrive in time to build the planes.

Decisions regarding whether or not to use titanium must reflect the

following realities:

0 Supply--today and 4 or 5 years ahead

0 Cost--measured by cost to save a pound of weight

0 Faci 1 iti es--forging, foundry, and machining capacities

0 Design problems--can other materials do the job for the same

weight?

P21-3
Supply

The key to the supply of titanium lies in sponge production. . When

Russia withdrew their supply, our ability to support ingot production became

tight. The supply of ingot material is very close to the market demand. By

importing some ingot and forgings, we are just meeting demand. It is our

speculation that the industry can supply our current and projected needs j_F

no new substantial demand is established. It is highly probable that

airlines will want more aircraft fuel efficiency. Our fuel-efficient

engines require titanium. We can apply titanium to build lighter airframes,

but if a military or other government program establishes priority

requirements on the existing titanium capacity, perhaps we should use

another approach to save weight, in order to protect our supply of engines.

Figure ^illustrates that nonaerospace products are beginning to take .

larger slice and that steadily increasing demand will probably continue

putting even more pressure on the supply part of the formula..

Cost

Currently, a 747 costs $70 million to an airline. We have to build and

test that airplane to their satisfaction for something less than that in

order to remain in business. We are today competing with whole groups of

nations, as well as other aerospace businesses, making our costs especially

sensitive in the decision process. Our way of deciding is to have our

various experts process the decisions on the basis of real cost to save a

pound of weight. Material selections must be made very early in the design

process, when little is known about part geometry, making those decisions

very difficult. Government titanium researach (AFML, U.S. Navy, U.S. Army,

NASA funded) has developed and offered technology to our industry, leading

P21-4
to lower, more competitive costs per pound. These activities must continue.

If, however, the material itself is prioritized out of existence to us in

the commercial world, the cost per pound evaluation becomes meaningless.

Faci lities

When we build an airplane at Boeing, we accomplish about 35% of the

total job. Our subcontractors and suppliers throughout the world do the

rest. Foundries, forge shops, machine shops, etc., and their production

capacities are an important part of any successful formula. In 1979, the

lead time to procure a large titanium forging was 155 weeks--nearly 3 years.

This is devastating when trying to support a production program to guarantee

delivery dates.

Design Problem

Some parts of an airplane have to be titanium. Engines, as mentioned,

must use substantial amounts of titanium. The basic fuel-efficient engines

are P&W-JT9, GE-CF6, and RR-RB211. Even our Airbus competitor uses these

engines and most, if not all, of the titanium in these engines comes from

U.S. capabilities (about 2,000 pounds per engine). In el evated-temperature

design applications (300 to 800°F), titanium performance cannot be

duplicated. In room-temperature applications, some of the higher strength

titanium alloy applications cannot be duplicated, but other parts can be

designed in steel, with only slightly higher weight. For subsonic aircraft,

most of the airframe can be, and is, designed best with aluminum alloys. An

increasing trend is to use graphite/epoxy composite hardware to replace

aluminum and fiberglass. Interestingly, this increases the amount of

titanium used for fittings and fasteners, since titanium is very compatible

P21-5
with graphite.

The design part of the formula works out in summary like this;

1. Engines must be designed in titanium.

2. Supersonic cruise air vehicles must be designed in titanium.

3. Certain parts of heavily loaded subsonic or transonic air vehicles

should be designed in titanium alloys.

4. Some parts of subsonic air vehicles and other products will perform better with titanium, depending on economic considerations.

To illustrate point 4, let's look at the 747. To replace only the

titanium parts that could be aluminum or steel would add approximately

10,000 pounds to the airplane. This means that to fly the long-range

segments of air routes for which most airlines use the 747, they would have

to take out over 50 seats or leave behind 10,000 pounds of revenue cargo.

The design part of the formula, in what we see of the future, weighs

heavily toward expanded use of titanium. There is, however, only a finite

amount of "fast footwork" available to our procurement people when the other

realities of the total equation are considered.

NEEDS

Paramount to Boeing is the need to establish a steady supply of the

various forms of titanium alloys as we need them. Our current commercial

airliner needs center on billet for forgings with smaller amounts of sheet,

strip, and bar. Increasingly, castings are becoming important, due in large

measure to the success of Air Force research and successful experience in

engines. Federal research in titanium needs to be continued in the

direction of higher strength alloys and lower cost and new ways of

fabrication. Powder metallurgy and superplastic forming/diffusion bonding

are fields where federal research has paid off, putting the U.S. well ahead

P21-6
of the rest of the world. Research in substitute metal systems which could

be used with no weight, cost or performance penalties needs to be

encouraged.

CONCLUSIONS

0 The availability of titanium is tight, but adequate, for today's

requirements.

0 Boeing will continue to expand its use of titanium and to use

technological advancements as they emerge.

0 A new supersonic cruise vehicle or major military production buildup

without a pound-f or-pound increase in titanium raw material capability

(primarily sponge capacity) would adversely affect our ability to

compete in the world commercial and military market.

0 Titanium technology, as led by federal research agencies and private

investments, has put this country well ahead of the rest of the world.

These activities need to be continued’.

0 Research in substitution technology needs to consider titanium.

P 21-7
RECOMMENDATIONS

0 Clearly establish that any expansion of military programs that will use

titanium requires an investment in titanium sponge production by

government or industry with tax credits or other monetary encouragement.

0 Bring stockpiles of titanium up to required levels by exercising a

planned expansion with purchases during industrial slumps and bartering

on international trade.

0 Insure that any research in substitution technology will consider

titanium and titanium alloys.

*

P21-8
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P21-9

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P21-10

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P21-11

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P21-12

FIGURE
TANTALUM AND OTHER CRITICAL ELECTRONIC MATERIALS

Thomas D. Schlabach

Bell Teleohone Laboratories

P22

KEYNOTE ADDRESS

Tantalum and Other Critical Electronic Materials

by

T. D. Schlabach

Bell Telephone Laboratories

Murray Hill, New Jersey 0797^

j_. Introduction

The electronics industry is a strong contributor to our

economy and to our technological leadership in the world.

In 1979, it had sales of $81 billion and a trade surplus of

nearly $5 billion while employing 1.5 million workers

according to the Electronics Industries Association and as

shown in Table 1. Its use of materials ranges from gross to

atom-by-atom additions in products that are very expensive

compared to most other products so that starting material

costs are seldom overiding. Nevertheless, the issues of

materials conservation and substitution are very important

to this industry and can present it with exceedingly difficult problems. These problems will be reviewed here for

several materials considered critical to our industry.

2. Electrical - Electronic Demand and Import Level

s

The criticality of an element to an indvistry includes

its level of use, its substitutability and the vulnerability

of ^.ts supply sources to interruption. Table 2 gives the

P22-1
2

percent demand in the electrical-electronic sector for 23 of

the most important elements to this industry and Table 3

shows their import levels. Those that combine a high level

of use with high import level include tantalum, gallium and

strontium though the latter has large and diversified

foreign sources. Those that combine moderate to high use

and import level with difficulty of substitution include the

platinum-group metals, tin, mercury and gold. The remaining

are less critical by these criteria though supply interruption of them can cause, and has, substantial problems.

Three elements on this list; namely cobalt, nickel and

tungsten, are dealt with elsewhere in this workshop and will

not be treated further here.

3 . Tantalum

Tantalum, by virtue of its large electronic use as seen

in Table 4 and its high import level, is a critical element

to the electronics industry and a strategic element to the

U.S. Imports of tantalum-containing concentrates and tin

slags provided 97 percent of U.S. apparent consumption which

amounted to nearly 1.5 million lbs. in 1980. Only 3 percent

of supply comes from recycling though secondary production

from run-around scrap is substantial. Worldwide resources

are judged adequate to meet projected world needs through

the year 2000 by the U.S. Bu Mines though highly dependent

on Zaire and Nigeria if U.S.S.R. resources are excluded.

The dependency of a good portion of tantalum supplies on tin

P22-2
3

slags, something like one-third, could portend future supply

problems because, while tantalum growth is estimated by the

U.S. Bu Mines at ^%/yr

.

out to the year 2000, tin growth is

estimated at only 1%/yr. High recent prices for tantalum

has encouraged supply expansion in Brazil, Canada and Australia .

4 . Capacitors

The principal electronics use for tantalum is in capacitors where it competes with the other types shown in Table

5 which lists U.S. sales for all except captive capacitor

production. Because tantalum solid electrolytic capacitors

exhibit high reliability, long life and high specific capacitance per gram combined with low noise and dissipation

factor, they have grown to command more than two-thirds of

the total electrolytic capacitor market (1). Aluminum electrolytic capacitors have made significant inroads, however,

in low impedance and consumer applications by technical

improvements and by their lower costs. At the same time,

notable improvements continue to be made in the capacitance

ratings of tantalum powders to increase their efficiency

through increased specific surface area. These ratings,

expressed as a capac i tance- vol tage product per gram, are

expected to increase from 5000 in 1980 to 10,000 in 1985 in

production capacitors. Laboratory samples have been

reported to exceed this value already and even higher values

are distinctly possible. The monolithic ceramic capacitors

P22-3
- -

are enjoying the fastest growth because of their use with

silicon integrated circuitry.

The substitution possibilities for tantalum capacitors

offered by aluminum el ec trol yt ics and monolithic ceramics

combined with continued improvements in capacitor-grade tantalum powder have decreased growth estimates in this area

from 7 percent to 4 percent. These changes together with

higher production capabilities being brought on stream by

higher tantalum prices improve the supply situation for tantalum over the near term. Beyond this, research has been

conducted on mobium, hafnium and titanium as possible substitutes in solid electrolytic capacitors without notable

success to date. This effort could be renewed if needed.

5 . Elements Presenting Difficult Substitution Problems

The elements shown in Table 6 are among those for which

substitution is very difficult, if not impossible, by virtue

of their unique properties and applications, and for two of

these we are heavily import dependent (2). Platinum-group

metals are essential for certain electrical contact uses and

indirectly essential in providing crucibles and other equipment for the growth of single crystals and glasses. For

} these uses there are no satisfactory substitutes. Mercury

is essential for wetted contacts and , again, no satisfactory substitutes exist though designs continue to be miniaturized and the stockpile contains nearly 3 years demand.

Tin is essential to solders although its use in electrical

P22-4
5

applications has decreased 30 percent over the past decade

due to lower tin contents in solders and the higher level of

interconnection on silicon chips and hybrid integrated circuits. Further, large holdings are in the stockpile and for

tins major use in containers viable substitutes exist that

could free a large fraction of tin demand.

Mercury is the only one of these three for which worldwide reserves appear inadequate to meet the world’s cumulative demand out to the year 2000. The fact that platinumgroup metals come principally from the Republic of South

Africa and the U.S.S.R. raises questions concerning our vulnerability to supply interruption.

6. Gold Conservation and Substitution

Gold presents a different problem in conservation and

substitution arising from its high cost and the fact that it

confers a degree of reliability in electronic devices and

apparatus that is unmatched. This led naturally to the use

of large amounts of gold in electronics and to intensive

efforts now to conserve and substitute for gold whenever

possible. Since most of this use is in the form of electrodeposited coatings, conservation steps have included selective area plating, thinner coatings and composite structures

with gold or gold alloy cladding. Beyond this, substitution

possibilities are being intensively examined for contacts

and some of these are listed in Table 7 where they are compared to gold. Other cheaper, precious metal alloys are one

P22-5
6

obvious possibility and base metal contacts are another.

Composites of the kind shown are a more radical departure

but being explored in certain less critical applications.

Even metallic conducting compounds have received attention.

In all of these cases, there is an attempt to match a substitute material with the particular requirements for that

application in the hope that it can prove adequate.

7 . Germanium

Germanium, gallium and indium are three elements critical to some of our newest electronics technologies and for

this reason it is approp'riate to examine their supply-demand

status. They are used currently in comparatively small

amounts and, while they do not pose strategic problems,

there is some concern about the adequacy of their short-term

suppl y

.

Germanium is obtained as a by-product concentrate from

zinc mining and potentially huge, but untapped, resources

exist in coal flyash where germanium concentrations of

several 100 ppm are common. In 1980, the apparent U.S. consumption of germanium was 77,000 lbs. with imports equal to

17 percent of this demand according to the U.S. Bu Mines.

The major use of germanium is in infra-red optical systems

which accounted for some 53 percent of uses in 1980 (see

Table 8). Under electronics, semiconductor use peaked in

1967 but germanium used as Ge 02 in optical fibers for

lightwave communications is an emerging application

P22-6
7

estimated to require 11,000 Ibs./yr. starting in 1981.

These fibers with germania-sil ica cores exhibit rvery low

optical signal attenuation and provide a low-cost, high

information-carrying capacity transmission medium. Interest

in lightwave communications is growing rapidly throughout

the world and will displace a certain fraction of coppers

use in communications. Because of substantial growth in

germanium usage in both optics and electronics, some shortage may appear in 1981 but there is every prospect of adequate supplies in 1982 and thereafter to meet the projected

growth of 5%/yr. through 1990.

Q. Gallium

Like germanium, gallium is obtained as a by-product in

the production of aluminum and zinc and is potentially

recoverable from coal flyash. Bauxite, the most important

source, contains 30-90 ppm gallium (3)* U. S. consumption

of gallium was reported at 9,500 kg in 1980 by the U.S. Bu

Mines with imports supplying 70 percent of this demand and

recycling none. Gallium’s predominant uses are electronic

(see Table 9) and require for the most part six nines purity

material. These uses include: 1) light-emitting diodes, 2)

other infra-red sources and detectors, 3) amplifiers and

oscillators in high-frequency systems, 4) bubble memory substrates, 5) photovoltaic devices, and 6) superconductors.

These and other gallium uses are anticipated to grow at

6%/yr. through 1990 but beyond that growth is highly depenP22-7
8

dent on the availability and large-scale commercialization

of gallium arsenide photovoltaic systems for the direct

conversion of sunlight to electricity. Far smaller, but

important emerging uses of gallium are as sources and detectors in lightwave communications systems and as bubble

memory substrates. These are both areas having very high

growth rates. It appears that there will be sufficient

world-wide production capacity and gallium resources to meet

anticipated world needs out to the year 2000.

Indium

Indium is obtained as a by-product of zinc production

and its supply is determined by this. World production of

indium was estimated at 1.6 million troy oz. in 1980 by the

U.S. Bu Mines and imports were 0.3 million troy oz. Modest

recycling occurs. U.S. production and consumption figures

are not available. Forty percent of indiums uses are in the

electric-electronic sector (see Table 10) and include:

solders, contact coatings, transparent and conducting coatings on glass, infra-red detectors and sources, and photovoltaic devices. As with gallium, the market having the

largest potential use is that of pho tovol taic s for energy

conversion. Infra-red sources and detectors, based on

InGaAsP, for lightwave systems and transparent conducting

coatings on glass are two important and growing areas.

While U.S. reserves to meet U.S. demand out to the year 2000

appear inadequate, world reserves to meet world needs appear

P22-8
9

adequate. In a number of its uses, indium has available

substitutes

.

1

0

. New Technologies and New Material U se Patterns

While it is vital to establish material conservation

and substitution possibil ities for our critical and strategic materials, none of these is likely to be more important than the new technologies developed to perform needed

functions in wholly novel ways. These so-called functional

substitutions (M) offer the greatest opportunity for profoundly altering, and often reducing, our use of materials

and energy though they are not solutions that can be rapidly

and simply implemented. The system-like nature of these

solutions require substantial time for success and implementation. Nonetheless, numerous information- and energyrelated technologies are underway, some of which are noted

here, that offer marked material and energy savings and that

can ease the problems facing us in this area.

P22-9
10

1

1

, References

1. Bernard, W. J. , "Developments in Electrolytic: Capacitors," Electrochem . Soc., 17^ : 403C-409C (1977).

2. Schlabach, T. D, , "Critical Materials and Substitution

Strategies in the Communications Industry," Metal Prog .

117(2): 57-62 (1980).

3. Rosi , F. D. , "A Survey of the Market, Supply, and Availability of Gallium," Mats . and Society £: 3^9-365

(1 980)

.

y. Chynoweth, A. G., "Electronic Materials: Functional

Substitutions:" Science 1 9

1

725-732 (1976).

P22-10
TABLE 1

U S. ELECTRONICS INDUSTRY. 1979

Factory Sales

U.S. Gov’t.

Exports

Imports

Employment

$80.6 billion (-M49b)

$20.2 billion (+ 14%)

$16.7 billion (+24%)

$11.7 billion (+10%)

1,464,000 (+11%)

EIA 1980

TABLE 2

ELECTRICAL-ELECTRONIC DEMAND.

PERCENT U.S. TOTAL DEMAND. 1980.

>90 Y, Ga

61-90 Sr, Pb, Ta, Cu, Be, Hg

26-50 Cd, Ge, Au, Se, Pd, Ag

10-25 Sb, Sn, Ni, Co, Al, Pt,

Zn, W, In

P22-11
TABLE 3

IMPORTS FOR ELECTRONIC ELEMENTS.

PERCENT APPARENT U.S. CONSUMPTION. 1980.

>90 Sr, Ta, Pt, Pd, Co, Bauxite

51-90 Ga, Cd, Sb, Sn, Ni, Zn, W

26-50 Hg, Au, Se, In

10-25 Cu, (Ge)

<10 Y, Pb. (Be), Ag

TABLE 4

TANTALUM. 1980

Mine Production: Canada, Brazil, Australia

Reserves: Zaire, Nigeria, Thailand

Electronic 66%

Machinery 26%

Transportation 8%

Substitutes: Capacitors, Al, Ceramics;

Superalloys H- Carbides, Nb

P22-12
TABLE 5

CAPACITORS. FACTORY SALES (MILLIONS) 1979.

Type Units $

Tantalum Electrolytic 899 262

Aluminum Electrolytic 205 165

Ceramic Dielectric 4,371 352

Film Dielectric 621 214

Other 60 36

Total 6,156 1,050

EIA 1980

TABLE 6

CRITICAL METALS, iflao

METAL CONSUMPTION. PERCENT GSA STOCKPILE,

elec. imp. aec. months demand

Pt-Qroup 24 87 16 9

Ho 52 49 21 31

Sn 15 65 25 46

P22-13
TABLE 7

CQMTACT PROPERTIES QF VARIOUS MATERIALS

MCTAL8 RISISTIVITY DENSITY COST

) mierahm-cm g/cm® * vol. ral Au

Au 2.3C 19.3 100

Au (plated hard) 4.6 17.5 90

Pd 10.8 12.1 19

Pd-40Ag 42.0 11.3 1

1

Sn 1 1.4 7.3 0.04

COMPOSITES

Ag-Epoxy 100 1.6 4

Qraphite-Polymer 10®-10^ 1.6 0.02

COMPOUNDS

RuOj 50 7.0 1.4

TABLE 8

GERMANIUM, 1980

Plant Production: W. Europe, U.S., U.S.S.R.

Reserves: U.S., Africa, W. Europe

Uses: Instruments + Optics 63%

Electronics 41%

Other 6%

Other Resources: Coal

Substitutes: Semiconductors, Si, Ga, In;

Optics, Zn, Se

P22-14
TABLE 9

GALLIUM. 1980

Production: U.S., Switzerland. Canada

Reserves: India, Surinam, U.S.

Electronics 93%

R-hD 7%

Dental Alloys <1%

Other Resources: Coal, Kaolin

Substitutes: Semiconductors, Si, Ge, In;

Displays, Liquid Crystals

TABLE 10

INDIUM. 1980

Refinery Production: Japan, U.S.S.R., Canada

Reserves: Canada, U.S., U.S.S.R.

Uses: Electronic 40%

Solders -f- Coatings 40%

Research + Other 20%

Substitutes: Semiconductors, Si, Ge, Ga;

Reactors, Hf; Solders, Bi

P22-15
TANTALUn IN HIGH TEMPERATURE ALLOYS

James A. Ford and Robert B. Herchenroeder

Cabot Corporation

P23
Tantalum in High-Temperature Alloys

James A. Ford and

Robert B. Herchenroeder

Cabot Corporation

Workshop on Conservation and Substitution

Technology for Critical Materials

June 15-17, 1981

Vanderbilt University

Nashville, Tennessee

P23-1
Tantalum in H igh~Temperature Alloys

INTRODUCTION

Robert B. Herchenroeder

In this presentation, three aspects of tantalum (Slide 1) will be considered:

(1) background, covering general information on uses of tantalum, past

Production, estimates of future production, and division of tantalum among

customers; (2) the main topic of the presentation which is the role of

tantalum in superalloys and the possibilities of substitution of other

elements; and (3) the possible "sleeping giant" consumer of this critical

material, the chemical process industry.

BACKGROUND

Slide 2 shows the sources of^ tantalum concentrates in the past three years and

indicates a near doubling of imports in this time period.

Slide 3 reviews the growth pattern of tantalum and tantalum mill products from

1976 through 1980 and gives an estimate of consumption of these products

through 1985. Note, the next five years appear to be a repeat of the past

five years.

Slide 4, similar to the previous one, but for tantalum remelt, predicts almost

a 70 percent increased use of tantalum for alloying in the period of 1980

through 1985.

Combined (Slide 5), the overall prediction is a drop in consumption of mill

products and remelt stock for 1981 and then a continual growth of about 50

percent through 1985.

Note this series of graphs does not include tantalum consumed for capacitors,

but it does include lead in wires associated with capacitors. The data are

for U.S. production and consumption only.

P23-2
During 1980, about 853,000 pounds of tantalum powder were used for

capacitors. Combining this with the other available data (Slide 6) it is

shown that only about 11.5 percent of the total consumed was used for alloying.

Slide 7 is a condensed listing of the applications for tantalum and its alloys

showing widely diversified consumers. Applications range from heat exchangers

to exotic electronic instruments.

Finally, in the way of background. Slide 8 lists a number of elements which

might be considered as substitutes for tantalum in superalloys. Since cost

also is indicative of availability, the list, indirectly, suggests how

suitable these elements would be for substitution in case of emergency. Costs

are for sponge, anodes, or high grade scrap.

P23-3
The Role of Tantalum in High-Tetnperature Alloys

Utilization of tantalum as a significant alloying agent is still in its

infancy. A review of the compositions listed in Slide 9 shows most

superalloys containing tantalum do so at levels of < 4 w/o. As indicated on Slide 9, it is this group of alloys which are currently being commercially

utilized. Despite much publicity about alloys containing 4%, i.e., as much as 12 w/o tantalum, these high tantalum alloys are still in the

experimental and developmental stages and do not currently constitute a significant drain on available supplies of tantalum; they may in the future.

Tantalum is used in commercial alloys of today (Slide 10) to improve strength,

surface integrity, and metallurgical stability. The penalties associated with

its utilization are mainly cost and increased density of components. However,

a $4-12/per pound increase in metal cost and perhaps a 10% penalty in

increased density is more than offset by as much as a doubling of rupture

strength (Slide 11),^ improved long-term ductility, and better corrosion

res is tance

.

The improvements in alloy properties (Slide 12), because of tantalum

additions, are related to solid solution strengthening, carbide stability,

carbide morphology, gamma prime stability, constitution and morphology,

improved castability, i.e., less freckling, and neutral or beneficial effects

on corrosion resistance compared to adverse contributions by other possible

alloying agents like hafnium, zirconium, and columbium.

Solid solution strengthening is achieved simply because the tantalum atom is

larger than that of the nickel (or cobalt or iron) base, and therefore,

strains the lattice when it is dissolved in the matrix. If this were the only

consideration, other elements such as molybdenum or tungsten could readily be

substituted

.

Carbide stability is a major reason and consideration for alloying with

tantalum. This is especially true for cobalt-base alloys such as 509 which

P23-4
contains about 4 w/ o tantalum. Tantalum monocarbides (TaC) tend to form as primary carbides during solidification of the molten alloy. With appropriate

solidification techniques, the form, size, and distribution of the TaC can be

controlled to yield spherical, script, or whisker formations.

In the case of alloy 509, the MC (M for metal because rarely are the carbides

pure TaC) form typically as Chinese script (Slide 13) and remain in that form even though the alloy is held at high temperatures for prolonged periods.

The NiTaC eutectic alloys are a superb example of advanced utilization of

TaC. These alloys containing 6-11 w/o tantalum, perhaps 0.04 w/ o carbon, plus

other typical alloying elements of cobalt, chromium, aluminum, titanium,

and/or vanadium, are slowly unid irec t iona 1 ly solidified in such a manner that

nominally 1 micron whisker is continuously grown along the longitudinal axis

2 of the casting yielding a natural fiber reinforced part. (Scredon gives a brief report of this.) Some of the characteristics of parts made from NiTaC

eutectics are listed on Slide 14.

It is reported that researchers have attempted to achieve similar results

utilizing other carbides than TaC, but with little success. Possibly, partial

replacement of tantalum in the NiTaC with another element could be successful,

but this is hypothetical.

As mentioned, the major reason for the use of tantalum as an MC former is its

stability. Slide 15 lists the predicted ranking of metal monocarbides based

upon the delta-free energy of the metal plus carbon reactions at 1400°K

3 (slightly lower than the temperature of formation of primary carbides).

Empirically it has been determined (also Slide 15) that both tantalum and

columbium monocarbides 'form in preference to TiC.

P23-5
Based on this intelligence, it would seem that columbium, a cheaper and more plentiful element, could logically be substituted for tantalum in these

h igh-temperature alloys, but this is not so. In practice, the evolution has

been from columbium to tantalum.

As mentioned, columbium carbide does not have an optimum morphology. Further, even though the simple thermodynamic calculations suggest TaC and CbC AG

R are close, i.e., -33610 and -32707 calories, respectively; other factors

influence the systems. Slide 16 illustrates these.

The refractory metal monocarbides are seldom pure. The MC , containing a preponderance of columbium, frequently and typically break down and dissolve

initially in the matrix. Depending upon the alloy system and temperature and

time lapse, the columbium may combine with the y' frequently displacing

chromium which then in turn combines with the available carbon to form

2366or M^C which usually precipitates in a detrimental manner. It may

combine with other elements of the alloy to form sigma. Laves, or other

topologically closed packed phases; but, almost without exception, the

columbium does not recombine to form CbC and the second state of the alloy is

less desirable than that of the first.

Hafnium, another possible substitution for

free-energy considerations, has been succes strengthened alloys for 10 to 15 years. Th

carbide and is used in amounts of 0.5 to 2%

boundary carbide precipitation. Typical al

which have enhanced ductility as a result o supplies of hafnium are as limited as those

for tantalum is not a good choice for that

tantalum, based on Gibb's

sfully used in gamma prime

is element does form a stable

to control the morphology of grain

loys are B-1900 + Hf and Mar-M-2A7

,

f the addition. Unfortunately,

of tantalum and substitution of it reason

.

With respect to tantalum's effect on y', it is generally accepted that the

element improves the h igh-temperature stability of the phase because it tends

to partition in like percentages in both the matrix and gamma prime. It does

not alter significantly the equilibrium of the matrix as does most other

components of the y

P23-6
With the advent of single crystal technology for gas turbine blades, tantalum

promises to be a major alloying element in the next generation of alloys.

Without grain boundaries, carbon, hafnium, and boron (grain boundary

strengthening elements) can be eliminated and tantalum, chromium, and tungsten

can be increased as has been done in PWA alloy 454 to improve strength and

oxidation resistance.

Copley, et. al.^ slide 17 explains that in addition to the others attributes of

tantalum, tantalum significantly reduces the tendency toward "freckling"* in

castings

.

Mr. Cell and co-workers at Pratt & Whitney Aircraft are reported to have said

"...this element (tantalum) alone among the third series transition elements

confers both high levels of creep strength and oxidation resistance."

Frequently, the substitution or replacement of tantalum for columbium in an alloy results in marked improvement of oxidation and hot corrosion

resistance. If one compares the oxidation resistance of WI 52 (a cobalt-base

alloy containing 2 w/ o Cb) with that of Mar-M-509 (a cobalt-base alloy

containing 4 w/ o Ta) this aspect is readily noted.

Slide 18 clearly shows that while tantalum may not improve oxidation

resistance pei se, it does not degrade resistance as does columbium. This is

from work done in the Technology laboratories during the development of HAYNES

alloy No. 556. The base composition was modified as indicated with varying

amounts of columbium, tantalum, and lanthanum.

* freckling in alloys results when elements such as chromium, and molybdenum segregate during casting to

separate bodies in the solidified alloy.

columbium, tungsten,

form distinctive.

P23-7
Fryburg, et. al. (Slide 18) noted that those alloys containing high

tantalum are more resistant to hot corrosion attack than those with lesser

amounts

.

A proposed mechanism of NaSO^- induced corrosion of B-1900 alloy is that the

protective ^l^O^ scale at 900°C is dissolved by liquid Na^MoO^ which

is formed by the reaction

MoO. (A) + Na„SO, (a) > Na„MoO, (SL) + SO- (g).

3 2 4 2 4 3 °

Once the Al20^ protective scale is penetrated, corrosion attack is swift.

When significant amounts of tantalum are present as in TRW VIA with 9 w/

o

tantalum, it is thought that the Na

2 S 0

^ reacts with fa 20 ^ to form

preferentially NaTaO^ , a solid, according to the reaction

Ta20^ (S) + Na2S0^ (£ > 2 NaTaO

2 .CS) + SO^ (g).

In this instance, the MoO^ volitilizes; the AI

2 O

2 protective scale

remains in tack.

Substitution of other elements in lieu of tantalum on the basis of oxidation

resistance and surface stability is not attractive. Columbium causes spallation of both Cr20^ and AI

2 O2 scales. Molybdenum forms liquid or gaseous MoO^ which can be most harmful especially in static atmospheres.

Zirconium, while beneficial in small quantities, reacts with mold materials,

collects along grain boundaries where it oxidizes preferentially, and forms

undesirable intermetallics when utilized in any substantial amount ("> .5

w/ o )

.

Tantalum is a unique element which beneficially affects many aspects of

superalloys. In general, substitution of one or more elements for tantalum in

the new generations of superalloys would incur meaningful penalties with

respect to almost all aspects of the resulting alloy's performance.

P23-8
In effect, the current and promised new era of cast superalloys would have to

revert to the technologies of the late 1960s and performance of equipment

utilizing these superalloys would be materially impaired.

%

P23-9
Tantalum Potential in CPI

Perhaps the largest eventual consumer of tantalum will be the chemical process

industry (CPI). As such, it can be a major factor in availability during a shortage of this potentially critical material. During 1981, a down year, it

is anticiapted that about 60,000 pounds of tantalum will be used in CPI. Much

of the future growth of mill products is expected to be in CPI applications.

From basic electrochemical determinations (Slide 20) one would expect that

tantalum, a relatively active material chemically, would have poor aqueous

corrosion resistance, but this is not the case in actual practice (see Slide

21). Tantalum and its "sister" element columbium tend to form thin amorphous

films that provide superb corrosion resistance through passivity to most acid

combinations

.

The near perfect resistance of tantalum to acids, as indicated, coupled with

the high cost of construction and maintenance of chemical plants, pleads well

the case of tantalum. Tantalum as a thin cladding on less corrosion-resistant

materials becomes more cost effective each day.

In considering critical needs along with possible shortages, one should

anticipate the likely high demands which will be associated with CPI.

i

P23-10
Summary

Tantalum is a unique alloying element for superalloys which affects numerous

aspects of an alloy's performance, including strength, surface integrity,

fatigue life, metallurgical stability, and castability.

The use of this element for alloying has been retarded and remains at a relatively low level because of cost and density penalties, but the advantages

achievable through new technologies such as directional solidification more

than counter these adverse effects. Use should increase dramatically. Cost

of tantalum in relation to total cost of an advanced design cored and

directionally solidified aircraft airfoil is insignificant.

Because of the special characteristics of this element for imparting

beneficial effects to alloys it is virtually impossible to substitute another

element and retain similar life expectancies from the parts made there from.

The chemical process industry will probably contribute to the shortage and

consumption of tantalum in the immediate future and this should be considered

in any plans for allocating this element.

P23-11
Re f erence s

1. N. N. Burova, et. al., "Structural Characteristics of Nickel Alloys with

Tantalum," Metal Science and Heat Treatment , May-June, 1979.

2. S. Scredon, "Next Generation Superalloy is Cast into Blade by G.E.,"

AMM-MN , August 18, 1980.

3. Barin, Knacke, and Kubaschewski , "Thermochemical Properties of Inorganic

Substances," Spr inger-Ver lag , 1977.

4. Atlas of Microstructures, ASM Handbook, 8th Ed., Vol. 7.

5. S. Copley, et. al., "The Origin of Freckles in Unidirectional ly Solidified

Castings," Me t . Trans

.

, Vol. 1, No. 8, 1970, p. 21.

6. G. C. Fryburg, et. al., "Mechanism of Beneficial Effect of Tantalum in Hot

Corrosion of Nickel-Base Superalloys," J. Elec. Chem. Soc., July, 1977.

P23-12
Slide 1

Tantalum in High-Temperature Alloys

Three Topics of Discussion

1. Uses and distribution of tantalum.

2. Role of tantalum in superalloys.

3. The "sleeping giant" consumer - CPI

RBHH

6/81

P23-13
Slide 2

I. SUPPLIES

u. s. Imports of Ta Concentrates. LbsxIO‘3

Country 1978 it ) 1979 it ) 1980 (X) Est

Australia 122 (9) 229 (15) 370 (15)

Brazil 179 (12) 179 (12) 550 (22)

Canada 667 (97) 679 (99) 573 (23)

China - - 102 (9)

Rwanda 123 (9) 131 (9) NA

Zaire 160 (11) 75 (5) NA

All Others 183 (13) 239 (16) 881

1929 1532 2976

U. S. Bureau of Mines Data

RBHH

6/81

P23-14
T.*‘J}TALUM

and

TAJJTALUII

ALLOY

Mill

Product

cu “O

on

o

to

09

S-ro

<1) >- O

a\

CO

\o

o o

^_0T X spunod

P23-15

Producer

Report

Data
TAi^TALUH

Hstorial

(Ta

additiona

for

high

tenperaturs

alloys)

I

1 . ^grrgaraae M&^uiitJBnMnnmBnw-ww ^

- 1

CO

n’i ^ gfcTwmj

|,.f

# iw ' 'Til* '1 ? 1

[ jv^v^rnsTOiV . wnsrvswnwv^

cr~

1 byrcn I'i. ^T.'rgrr.'^T ^v.wwnar

i

tKfr»Td3CTanfc iT»rs5pw.xr*7St7T:

1

1

1 «r^TSTk-nu rr‘)trj-v.%rr^'^:^n:'‘ 5 5

; —• irsivffrrri'.rnc'i'CB’^niMiw

1

»

AT.i'yr .-laf jJiVT- iwt* z?aan®*t

. PTO P'<-/38: wWMMrtifesr'rwwJ*-

:

or 3 —r

o o o o o

C3 o o o o m fO OJ r~) ^_0l X spunod

P23-16

tno

r?03

fO

CO

C'l 03

C3

O

CO

as

p'

CD

so

S-n3 OJ >-

Producer

Report

Data
TOTAL

TANTALUM

(Mill

Products

+

Ta

Rcrr.olt

QJ “D

i 1

O

o

o

CD

r>

CO

(NCO

CO

S-ro

<V >-

O

rCO

r'

KO

o

V? o

'3C' o

(N

g_ 0 I X spunod

P23-17

Producer

Report

Data
Slide 6

1000 Lbs. Percent

Tantalum and Mill Products 365 26.5

Tantalum Remelt 159 11.5

Capacitor Powders 853 61.9

TOTAL 1377 99.9

NOTE: Assuming 2.5 w/o Ta content for the average content

OF A TANTALUM-CONTAINING SUPERALLOY, THIS MEANS 6,360,000 POUNDS

WERE PRODUCED IN 1980.

RBHH

6/81
Slide 7

TANTALUM AND TANTALUM ALLOY PRODUCTS

Pure Metal: Equipment for Chemical Industry

Heat Exchangers

Condensers

Heaters

Distallation Columns

Reaction Vessels

Hardware

Electrical & Electronics

Capacitors

Halogen Lamps

Infra-Red Lamps

Power Sources

Lithium Battery Seals

X-Ray Hardware

Cathodic Protection Systems

Vapor Deposition

High Temperature Furnaces

P23-19

Heating Elements

Shields

Support Hardware
Slide 7 (Cont. )

Ta-2-1/2%W Alloy Equipment for Chemical Industry

Same Applications as Pure Metal

Electrical & Electronics

X-Ray Hardware

High Temperature Furnaces

Same as Pure Metal

Ta-90%W Alloy High Temperature Furnaces

Same as Pure Metal

Aerospace

RBHH

6/81

-f

P23-20
Slide 8

Relative Cost of Materials*

Dollars/Pound

Tantalum

Hafnium

COLUMBIUM

Zirconium

Titanium

Vanadium

85

85

25

12

7.25

4.60

• Sponge or high grade scrap

RBHH

6/81

P23-2I
SUPERALLOYS

CONTAINING

TANTALUM

X

M cu 0) C|- X q: x: X lA lA

• • O r—

1

o

m| m|o o

1—1 1—

i

1

<

1 r—

(

MD 05 60 05 *—H f—

H

.13 O O o o o O o o o

o

CQl

.015 .015 CN 1 CM

1 O O O

IPi .-1 CM r-H

1 O CD CD

CD CD

CD CD .02 o o CD CD CD CD CD CD CD CD CD

CN r- CM rA LA MD lA CN CN rA

CJl

•—H ^

'—1

•—

1

1—

1

rH i—j rH r-H

• • 1 • • • . • • • • • o o CD CD CD o U CD CD o o o

LA <r LA LA o rA o

•Hi •—1 • • • • • 1 1 t-| •—

(

rA rH <3* I—

1

LA <r rH rA LT\ tn CO <r

VO ND LA • • • • ... VO VO • cel VO rA rA VO m lA CM LA

Xll Ch1 1 1 ^ • 1 1 1 1 CM

LA LA LA ol o CN rH O

cd| 4 4 2 (N • Ov o LA

• rA CN CD 03 ov

i-l rH rH rA rA rH

si

VO CD

1 1 CM • • o

o o rA <3* 00

CM rA 11 1—1 rH LA

(J\ 0 LA VO o VO VO 1 • • • . •

1 CM

2; i-H rH CN CM 0

c

LA LA 0

0 0 CD LA 1 • CJ\ 1 0 0 0 1 LA . •H

CJ .—

1

rH 00 rH CM 4-> CDU

•H

<r CN LA rH H CO 00 0 CD VO • • ov • . VO VO . CL 0 rH rH rH (N LA CO LA VO CL

rH •—

1

CD rH

CD •H rA CN rA rH rH <—

t

0 0 0 CD <r 1—

1

•H

VO VO VO VO VO VO VO LA VO VO VO U

M

CUE

E

c 0

0 U

•H

4-) ct -»-> CD (U IH CO C in C|- r-' CD > 0) CJ CD X VO CM a sz •H CD .-N + CN LA v—. ct CJ) CO I *-> 0 0 <r 00 CN CM 1 <r CD •H

0) (—H CO 0 0 LA CN rA (JN 1 X 1 ct CD ct X Q (U CD CTv CJn < VO 1^ 5; 1 CO CD -z. U t »—

1

rH 1 cr 1 1 > 0 ^ 1 1 cl X X CN cc ct DC M M s * 0 •H CD CQ X t-H HH CN ct X ct ct cc X rH

rH * CL t-H 5; •Z. * 2: 1— 1— 1— X 00 rH CD ct P23- 22

q: VO
SUPERALLOYS

CONTAINING

TANTALUM

r

C -J X mo o o

0) •H

0) Li_ to X IT\ X • • o r-H o

CO _iCNOO

(/)

O) U- o

(-il CM lA

C f—1

M| • • X • o o o LA o

i—H

lA CU O U- •—1

col CD 1 o

«—1 O CO o

CD

LA LA

VO CO VO 1—

H

rA

cjI o O

O o o

<N

*r-| 1 • CM o • rA o

<N rA VO

f-H 1 • * * O ' o

r-H -qI 1 1 • 1 cjI * o

CO LA cjv LA

rol • Ch • • • LA o r—j

LA s| o r' • VO '—1

CN

O 1 1 1 rA rA

2;

CO LA CO o o CO LA LA •—H

CJ CD

LA

M CN CM • CN CN CN CM LA 4-J CN r-H X

crDO i VO 1 O J-i

O C7N 2 1—1

CN lA

<D

(/) VO

CO LA

C CD LA o

•H c

4-) o O

CO (U CN I—< c (/) O CA HH r—j o CO O CM O CO

•r- CD ^ 1—

O LA "O

(/) 1 -L-> LA 1 c CO <C

d) 4-J (/) CO t 2 CO LjJ

Q

Q CO UJ s; 1 X h-i CO C_) X 1 cc r—4

>- 1 > X ' >- cr cr (D Ct CN O o cl CC 2 X U.

»— CJ

X 2 * U ct

r— •rH P23-23 co

•H

CO

L)

Ql

CL

CO

CO

•rHO

0)EOa

</) (D JCO) X

X 00

CQ \

rr vn
Slide in

Benefits and Penalties of

Tantalu m in Superali nyt;

Benefits:

Improved strength

(perhaps double)

Improved surface integrity

(.5 TO .1 CORROSION RATE)

Enhanced metallurgical stability

(70%+ RETAINED DUCTILITY)

Penalties

:

Cost (?q-i2/LB. of metal)

Increased density (up to 10%)

RBHH

6/81

P23-24
CORROSION

RATES

(MILS

PER

YEAR)

6/81
roO

o CO ^ ^ ro <NJ

CO

LO

CO

CM

C\Jo

1“^ o>

CO

LT)

ro

CM

uoO

iuui/^6>| ‘ss0j:^s

c

P23-26

After

N.

N.

Burova

and

S.

B.

Maslenkov

^
Slide 12

Reasons for Tantalum Benefits

Solid Solution Strengthening

Carbide Stability

Carbide Morphology

Gamma Prime Stability Constitution and

Morphology

Less Freckling

Beneficial or Neutral Effects on Corrosion

RBHH

6/81

P23-27
Slide 13

Kallinfi's reagent lOOX Electrolytic 6''', ptiosphonc ncid 500- 1590 MAR-M 503, as cast. The stiaictiirc eon- 1591 Same a,s 1591. but at a lil.alicr maunlsist-s of MC iiai'ticlcs in .seiipt fni in i- ^ -J i- (nation, to reveal niorpliology of MC

and M. C.. particle.s in eutectic form (gray areas) serijR particles, ijriinary eutectic particles and precl|)itate form in tlie dendritic alplia (M.C.i. and |)reclpilated M: C.. (shadowy consolld-solution matrix (facc-ccntcrcd cubic). stituent' in tire alpha (fee) matrix.

P23-28
Slide 1M

NiTaC Eutectics ^

(6-11 Ta .OMC + Co, Cr, Al, V)

1. Oxidation resistance comparable to R80.

2. Operation ^ 250°F higher temperature results

IN 1% INCREASE IN FUEL COUPLED WITH 17X INCREASE

IN THRUST.

3. Nominally 1 micron whisker oriented along

LONGITUDINAL AXIS.

M. Reportedly as much as 20 longer life in low

CYCLE FATIGUE.

5. 30% HIGHER ENDURANCE STRESS AT HIGH CYCLE FATIGUE.

6. Researchers are reported to be evaluating

Cb, Cr. and Mo without appreciable success.

RBHH

6/81

P23-29
Slide 15

Thermodynamic Prediction of Stability

- Cal.

Hf ^ C - HfC -52723

Zr ^ C - ZrC -H3895

Tl > C

-* TiC -H0088

Ta ^ C - TaC -33610

Nb * C - NbC -32707

V ^ C - VC -21335

w + c WC -8310

Empirically order OF FORMATION

RBHH

6/81

HfC, ZrC, TaC, CbC, TiC, VC. WC

P23-30
Slide 16

MC TYPICALLY NOT PURE, I.E., TiC, TaC, ETC., WITH OTHER "M" ATOMS SUBSTITUTING

FOR ONE ANOTHER.

EXAMPLE:

(Tao.8 Moo.2)C

N 13AL OR o OR Laves ^23^6* MeC

RBHH

6/81

P23-31
Slide 17

Improved Castability^

Steve Copley, et. al., reports tantalum significantly

REDUCES TENDENCY OF "FRECKLING" IN CASTINGS.

Metallurgical Transaction, Vol. 1. No. 8, 1970.

RBHH

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P23-32
14

Slide U * g

COE

<Vu

H3

CO

CO

350

300

250

200

150

100

51

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

CoTumblum Content, Weight Percent

DYNAMIC OXIDATION DATA VS COLUMBIUM CONTENT

P23-33

Metal

Loss/Surface

,

microns
Slide 19

Tantalum Influence on Hot Corrosion^

Case I - No Ta

M0O3 + NA2S0q (^) ^ NA2Mo0q + SO3 (G)

NA2M0O4 DISSOLVED AL2O3 SCALE

Substrate corrodes severely

Case II - 9 Ta

M0O3 + TA2O5 (S) + NA2SO4 (^)

- 2NATAO3 (S) + SO3 (6) + M0O3 (G)

NATAO3 IS SOLID AND DOES NOT FLUX AL2O3

Substrate remains protected by AL2O3

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P23-34
Slide 20

ELECTROCHEMISTRY

Std. Oxid. Pot.

Cb Cb^^^ + 3e" +1.10V

Ta TA+++++ + 5e- +0.85V

Cb + 5e" +0.65V

Nl Nl++ + 2e" +0.25V

Mo Mo^^^ + 3e" +0.20V

Pd Pd'^'" + 2e~ -0.98V

Au Au^^^ + 3 e- -1.50V

Active

Less Active

Noble

In theory Ta is quite active!

It is presumably the special properties of "amorphous" film

THAT PROVIDES EXCELLENT CORROSION RESISTANCE THROUGH PASSIVITY.

RBHH

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P23-35
Slide 22

TRADEMARKS

MAR-M Martin Marietta Corporation

HAYNES Cabot Corporation

HASTELLOY Cabot Corporation

MULTIMET Cabot Corporation

INCONEL INCO FAMILY OF COMPANIES

INCOLOY INCO FAMILY OF COMPANIES

RENE' General Electric Company

UDIMET Special Metals Corporation

UNITEMP Universal Cyclops Corporation

RBHH

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P23-36
Partitioning

in

TRW

VIA

alloy

P23-37

6/81
NBS-IUA (REV. 2-80

U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO. NBSIR 82-2495

2. Performing Organ. Report No. 3. Publication Date

April 1982

4. TITLE AND SUBTITLE

conservation and substitution technology for critical MATERIALS--Proceedings of

Public Workshop sponsored by U.S. Department of Commerce/National Bureau of

Standards and U.S. Department of Interior/Bureau of Mines. VOLUME I.

5. AUTHOR(S)

Alien G. Gray

6. PERFORMING ORGANIZATION (If joint or other than NBS. see instructions) 7. Contract/Grant No. national bureau of STANDARDS

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

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

Department of Commerce Department of the Interior

National Bureau of Standards rind Bureau of Mines

Washington, DC 20234 Washington, DC 20241

10. SUPPLEMENTARY NOTES

Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11. ABSTRACT (A 200-word or less factual summary of most significant in formation. If document includes a significant

bibliography or literature survey, mention it here) The United States is highly vulnerable to problems in supply of critical and

strategic materials and it is recognized that there is a whole spectrum of options

for responding to such crises. While a number of supply oriented options are under study by various groups, the focus of this Workshop was on the technical

opti ons

.

The Workshop was held orincipally to develop information for the report

required by the Department of Commerce, but should also be useful to the ether

agencies in their responsibilities. The DoC report is supposed to identify a materials needs case related to national security, econoniic well-being, and

industrial product! vi ty , to assess critical materials needs, and to recommend

programs to meet these needs.

12. KEY WORDS (S/x to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolon s) Critical materials; strategic materials; conservation; substitution;

chromium; cobalt; titanium; tantalum.

13. AVAILABILITY 14. NO. OF

PRINTED PAGES

[j^ Uni imited

1 1 For Official Distribution. Do Not Release to NTIS 488

1 1 Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402. 15. Price

Order From National Technical Information Service (NTIS), Springfield, VA. 22161 $ 34.50

USCOMM-DC 6043-P80

*U.S. GOVERNMENT PRINTING OFFICE: 1982-0-360-997/2137