NBSIR 83-2679-2 (£)

Technical Aspects of

Critical Materials Use

by the Steel Industry

Volume II B

IIProceedings of a Public Workshop; “Trends in Critical Materials

Requirements for Steels of the Future; Conservation and

Substitution Technology for Chromium”.

4-7 October 1982

Vanderbilt University

Nashville, TN

Public Workshop Sponsored by:

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

U.S. Department of the Interior, Bureau of Mines

U.S. Department of Defense, Army Research Office

June 1983

Center for Materials Science

U.S. Department of Commerce

National Bureau of Standards

GC —100

,U5G

33-2679-

1933

Z

RATIONAL BUREAU OF STANDARDS

LIBRARY

NBSIR 83-2679-2

Technical Aspects of

Critical Materials Use

by the Steel Industry

Volume II B:

Proceedings of a Public Workshop; ‘Trends in Critical Materials

Requirement for Steels of the Future; Conservation and

Substitution Technology for Chromiun.”

Sponsored by:

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

U.S. Department of the Interior, Bureau of Mines

U.S. Department of Defense, Army Research Office

4-7 October 1982

Vanderbilt University

Nashville, TN

Chairman

Allen G. Gray

American Society for Metals

Adjunct Professor of Metallurgy

Vanderbilt University

June 1983

U.S. Department of Commerce, Malcolm Baldrige, Secretary

National Bureau of Standards, Ernest Ambler, Director
Contents

VOLUME I: SUMMARY REPORT

EXECUTIVE SUMMARY ix

INTRODUCTION 1

PROCESSING 6

Current Economic Status of the Industry 6

Technical Status of the Three Major Sectors ... 7

Technical Developments in Processing 9

Continuous Casting 9

Primary Melting and Refining Processes ... 10

Real-Time Control 13

Near Net Shape Technology 15

Surface Modification 16

Steel Plant Refractories 18

Import Implications of Advanced Technical

Developments 19

Recycling 20

SUBSTITUTION 23

Chromium Functions in Steel 24

Substitution Options for Chromium in Corrosion

and Oxidation Resistant Stainless Steels ... 24

Substitution Options for Chromium in Structural

Alloy Steels 29

INSTITUTIONAL FACTORS 32

Specifications 32

Qualification 32

Composition of the Natural Defense Stockpile . . 33

Research 34

CRITICAL MATERIALS AND PRODUCTIVITY 37

REFERENCES 39
VOLUME II A: Proceedings of a Public Workshop; "Trends in Critical Materials

Requirements for Steels of the Future; Conservation and

Substitution Technology for Chromium".

Significance of the Workshop

CRITICAL MATERIALS NEEDS FOR FUTURE STEELS AND THE CHALLENGE OF THE

CHROMIUM SITUATION, Allen G. Gray, Workshop Organizer and Chairman;

Vanderbilt University and American Society for Metals. PI

Workshop Keynote Session

INDUSTRIAL LIFE WITHOUT CHROMIUM-TECHNOLOGICAL CHALLENGES, Arden L. Bement, Jr., TRW Inc. P2

POTENTIAL FOR CRITICAL MATERIALS CONSERVATION IN THE INTEGRATED STEEL

INDUSTRY, Gordon H. Geiger, Chase Manhattan Bank. P3

POTENTIAL AREAS FOR CHROMIUM CONSERVATION IN STAINLESS STEELS,

R. A. Lula, Consultant. P4

THE VALUE OF GENERIC TECHNOLOGY: SUBSTITUTION BASED ON

HARDENABILITY, Dale H. Breen, Gear Research Institute. P5

Users' Views on Steels Needed For The Future Technological

Trends on Critical Materials Required For These Steels

TRENDS AND NEEDS FOR FUTURE STEELS IN BUILDINGS AND BRIDGES,

Lewis Brunner, American Institute of Steel Construction. P6

THE MATERIALS USE AND RESEARCH OUTLOOK IN THE RAILWAY INDUSTRY,

William J. Harris, Association of American Railroads. P7

TRENDS AND NEEDS FOR FUTURE AUTOMOTIVE STEELS, George H. Robinson,

General Motors Corporation. P8

CATERPILLAR TRACTOR CO. VIEWS ON STEELS NEEDED FOR FUTURE

APPLICATIONS, Dennis B. O'Neil, Caterpillar Tractor Co. P9

CHEMICAL EQUIPMENT - USER'S VIEWS ON STEELS NEEDED FOR THE FUTURE,

Edward A. Kachik, Materials Technology Institute. P10

ELECTRIC UTILITY VIEW OF THE USE OF CRITICAL METALS IN STEEL,

Robert I. Jaffee, Electric Power Research Institute. Pll

THE USE OF CHROMIUM IN STEELS FOR AEROSPACE, Rod Simenz,

Lockheed-Cal ifornia Co. P12

USER'S VIEWS - OIL COUNTRY USAGE TRENDS IN CRITICAL MATERIALS FOR

STEELS OF THE FUTURE - CHROMIUM, John W. Kochera, Shell Development

Co. P13
Impact of Developments in Manufacturing and Process

Controls on Conservation and Recovery of Critical Materials

STAINLESS AND SPECIALTY STEELS AOD, EBR, LR, VAR, VIM, VOD, James

T. Cordy, Universal Cyclops Specialty Steels Co. P14

MINI MILLS-TECHNOLOGICAL INNOVATIONS AND FUTURE ALTERNATIVES,

Peter H. Wright, Chaparral Steel Co. P15

FUTURE RAW-MATERIAL REQUIREMENTS FOR STEEL PLANT REFRACTORIES,

David H. Hubble and K. K. Kappmeyer, U.S. Steel Corp. P16

RECYCLING-PRESENT AND FUTURE: "POTENTIAL FOR CONSERVATION"

Herschel Cutler, Institute of Scrap Iron and Steel. P17

Conservation and Substitution For Chromium in Stainless

Steels for Chemical Use and For Corrosion Resistant Applications

AN OVERVIEW OF THE POTENTIAL FOR CHROMIUM CONSERVATION IN STAINLESS

STEELS FOR CORROSION APPLICATIONS, Gerald L. Houze Jr., Allegheny

Ludlum Steel Corporation. P18

SUMMARY OF STUDIES OF METAL PROPERTIES COUNCIL ON CHROMIUM

CONSERVATION IN STAINLESS STEELS FOR CORROSION APPLICATIONS,

A. 0. Schaefer, Metal Properties Council. P19

CHROMIUM CONSERVATIONS IN STAINLESS STEELS: STATUS OF MPC ACTIVITY,

Jim Heger, Consultant, Metal Properties Council. P20

CHROMIUM SUBSTITUTION PLANNING AT THE AMERICAN STERILIZER COMPANY,

Roy S. Klein, The American Sterilizer Co. P21

ALTERNATES FOR STAINLESS STEELS IN THE CHEMICAL PROCESS INDUSTRIES,

Edward A. Kachik, Materials Technology Institute of Chemical

Processing Industries. P22

OPPORTUNITIES FOR CONSERVATION OF CHROMIUM IN CHEMICAL PROCESS

EQUIPMENT, Robert A. Gaugh, ARMCO Inc. P23

APPROACHES TO CHROMIUM CONSERVATION IN MATERIALS FOR CHEMICAL

PROCESSING INDUSTRIES, Aziz I. Ashpahani , Cabot Corporation. P24

POTENTIODYNAMIC CORROSION BEHAVIOR FOR SEVERAL FE-MN-AL AUSTENITIC

STEELS, Rosie Wang and R. A. Rapp, Ohio State University. P25
VOLUME II B: Proceedings of a Public Workshop; "Trends in Critical Materials

Requirements for Steels of the Future; Conservation and

Substitution Technology for Chromium".

Conservation and Substitution For Chromium in

Stainless Steels and Alloys For Heat Resistant Application

OUTLOOK FOR CONSERVATION OF CHROMIUM IN SUPERALLOYS, John K. Tien,

Juan M. Sanchez, and Robert N. Jarrett, Center for Strategic

Materials, Columbia University. P26

DEVELOPMENT OF 9 CR-1 MO STEEL, Vinod K. Sikka, Oak Ridge National

Laboratory. P27

S IL ICON-MOLYBDEN IUM DUCTILE IRON FOR ELEVATED TEMPERATURE SERVICE

TO CONSERVE CHROMIUM, Jan Janowak, Climax Molybdenum Co. P28

THE 1982 STATUS REPORT OF MN-AL-FE STEELS AS REPLACEMENT FOR

STAINLESS IN HEAT RESISTING AND CRYOGENIC APPLICATIONS,

Samir K. Banerji, Foote Minerals Co. P29

Conservation and Substitution For Chromium

in Carburizing, Heat Treatable Steels and Bearing Steels

THE DEVELOPMENT OF NEW ALLOYS TO REPLACE CHROMIUM IN CARBURIZING

STEELS FOR GEARS AND SHAFTS, Carl J. Keith and V. K. Sharma,

International Harvester Co. P30

CHROMIUM-FREE STEELS FOR CARBURIZING, George T. Eldis, D. E. Diesburg, and H. N. Lander Climax Molybdenum Co. P31

RARE EARTH BORON STEEL (25-Mn-Ti-B) GEARS REPORT FROM NANCHANG

GEAR PLANT, BEIJING, CHINA, Presented by Dale H. Breen and Allen G.

Gray. P32

POTENTIAL FOR SELECTIVE HARDENING BY INDUCTION IN CHROMIUM-FREE

STEELS, Peter A. Hassell, Ajax Magnethermic Corp. P33

BEARING^STEELS OF THE 52100 TYPE WITH REDUCED CHROMIUM, Chester F. Jatczak, The Timken Co. P34

Conservation and Substitution For Chromium in Structural

Alloy, High Strength, and High Strength Low Alloy Steels

AN OVERVIEW OF CONSERVATION AND SUBSTITUTION FOR CHROMIUM IN

STRUCTURAL ALLOY, HSLA AND ULTRA HIGH STRENGTH STEELS, Robert T.

Ault, Republic Steel Corporation. P35

ALTERNATIVE COMPOSITIONS FOR FUTURE HSLA STEELS, Brian L. Jones,

Niobium Products Co. P36
Potential For Advanced Technologies in Chromium

Conservation-Coating Systems And Surface Modification

Technology; Ceramics, Composites And Intermetal 1 ics

OPPORTUNITIES FOR SURFACE MODIFICATION TECHNOLOGY IN CONSERVATION OF

CHROMIUM, Peter G. Moore, Naval Research Laboratory. P37

SALT BATH TREATING AS AN ALTERNATIVE FOR CHROMIUM PLATING,

William G. Wood, Kolene Co. P38

CLAD METALS: MATERIAL CONSERVATION THROUGH DESIGN FOR CORROSION

CONTROL AND HIGH PERFORMANCE, James T. Skelly, Texas Instruments,

Inc. P39

POTENTIAL FOR POLYMER CONCRETE TO CONSERVE ALLOYS IN ENGINEERING

APPLICATIONS, Jack J. Fontana, Brookhaven National Laboratory. P40

ELECTROLESS NICKEL AS A SUBSTITUTE FOR CHROMIUM PLATING IN

INDUSTRIAL APPLICATIONS, Ronald N. Duncan, Elnic Inc. P41

DEVELOPMENT OF DUCTILE POLYCRYSTALLINE NKAL FOR HIGH TEMPERATURE

APPLICATIONS, C.T. Liu and C.C. Koch, Oak

J Ridge National Laboratory. P42

INJECTION MOLDING CERAMIC PARTS FOR HIGH TEMPERATURE APPLICATIONS,

Beebhas C. Mutsuddy and Dinesh K. Shetty, Battel le-Columbus

Laboratories . P43

Information Stockpile - Summary of Comments

DEVELOPING AN INFORMATION STOCKPILE TO AID IN SUBSTITUTION

PREPAREDNESS, Robert T. Nash, Vanderbilt University. P44
OUTLOOK FOR CONSERVATION OF CHROMIUM IN SUPERALLOYS

John K. Tien, Juan M. Sanchez and Robert N. Jarrett

Center for Strategic Materials

Henry Krumb School of Mines

Columbia University

New York, New York 10027

I. INTRODUCTION AND BACKGROUND

I Chromium is widely recognized as an essential alloying

element in superalloys. This point is illustrated in Table I,

which shows chromium contents for typical superalloys. Wrought

nickel-base and iron/nickel-base alloys constitute the majority

of superalloys and contain roughly 15-20% Cr. Cast nickel-base

alloys contain typically 10-15% Cr and cobalt-base alloys 20-30%

Cr.

Since the total tonnage of superalloy production is a mere

fraction of the steel production, the amount of chromium used in

superalloys does not constitute a very sizable portion of the

12 3 total end use chromium in the U.S. ' ' (see Fig. 1). However, the

readily available ferrochrome used in stainless steels cannot be

used as the source of chromium in the production of superalloys,

since steel-grade ferrochrome usually contains too many

deleterious tramp elements and impurities. Moreover, iron

certainly is a major contaminant in the nickel-base and

cobalt-base superalloys. Thus, the more expensive vacuum grade.

P26-1
2

TABLE I. Chromium Content of Typical Superalloys

Wrought Ni-Base Cr Content, %

Waspaloy 19.5

Udimet 720 18.0

Udimet 700 15.0

Nimonic 115 15.0

Cast Ni-Base

IN 738 16.0

Mar M200 9.0

IN 100 10.0

Alloy 713C 12.5

Iron-Base

Inconel 718 18.5

Alloy 901 13.5

Incoloy 903 0.0

Cobalt-Base

X 40 25.5

Mar M509 21.5

Mar M322 21.5

Waspaloy, Udimet, Mar M, and Nimonic, IN, Inconel and Incoloy are

trademarks of United Technologies, Inc., Special Metals Corp.

,

Martin-Marietta Inc., and International Nickel Company, Inc.,

respectively.

P26-2
3

spectral grade, or electrolytic* chromium metal is required.

Figure 2 shows the superalloy use of refined chromium and since

over 50% of this chromium metal is refined overseas, this data

clearly indicates the critical role played by chromium in

superalloy production. Although superalloys use only a small

fraction of the total chromium imported, they are the largest

user of high quality chromium metal for which the U.S. has a limited production capacity.

In order to establish the feasibility of a Cr conservation

program in superalloys, we will next briefly assess the role

played by chromium in the structural (or bulk) properties and in

the corrosion (or surface) properties of these alloys.

As a bulk alloying element, Cr is found to be the major

component of ^23^6 and carbides, a minor component

2? -i 7 ft of MgC, and not present in MC carbides ' ' (Table II). In

cobalt-base alloys carbides play a very important role since they

are the strengthening phases. Table III shows the basic types of

cobalt-base alloys. Note that the total amount of carbide

formers in these alloys are about constant, but as the tungsten

to chromium ratio increases, M

g

C carbides replace the

Cr 23^6* Furthermore, with the addition of Ti-Ta type

* Difficulties also arise in using electrolytic chromium due

to salt impurities resulting from that process (4).

P26-3
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TABLE II. Carbides in Cobalt-Base Superalloys

Alloy C Cr W+Mo Ti+Cb+Ta+Zr

Total

Carbide

Formers

Carbide

Type

X 40 0.5 25.5 7.5 - 33.0 Cr 23C

6

Mar M509 0.6 21.5 7.0 4.2 32.5

<W?3rLc

(Ta,TiJC

Mar M322 1.0 21.5 9.0 7.5 38.0 Wc C

TaC

TABLE III. Carbide Types in Superalloys and Their Metal Content

Carbide Metal Atoms

m

7c

3 Cr

M

23C6 Cr + (Mo f W)

M

6 C (Mo r W) + Cr

MC (Hf ,Ta,Nb,Ti,V)

P26-4
5

refractory elements, the Cr 23Cg carbides are replaced with

MC's which do not contain Cr. In fact, refractory metals are

purposely added to Co-base superalloys in order to replace

chromium carbides, which in general play a detrimental role in

the alloy's mechanical properties. When Cr 23Cg forms, it

precipitates heterogeneously at the grain boundaries and at

stacking faults. This stacking fault precipitation, in turn,

acts as a crack initiation site resulting in low temperature

embrittlement. Chromium, however, is always present in Co-based

superalloys at a concentration of about 20% Cr, in order to form

a protective oxide layer. The latter is necessary since cobalt

oxide is volatile, offering no protection of its own.

The same carbide decomposition reaction seen in Co-base

superalloys is used in the heat treatment of nickel-base

superalloys. However, in the latter system, the reaction is used

to produce the typical grain boundary morphology shown in Fig. 3,

with both forms of carbide (MC and M^Cg) present after heat

treatment.

In a study* on the role of cobalt in superalloys, we found

that cobalt affected the type of carbide found in nickel-base

superalloys^'^ (see Fig. 4). Removing cobalt from the standard

"The Role of Cobalt on Nickel-Base Superalloys", NASA COSAM

(Conservation of Strategic Aerospace Materials) Grant NAG

3-57 conducted at NASA-Lewis Research Center, Special Metals

Corporation, Purdue University and Columbia University.

P26-5
6

alloy destabilizes the MC carbide and results in more

Despite this effect the mechanical properties were essentially

unaffected (Fig. 5), indicating again that the chromium carbides

are replaceable.

As part of the same investigation, it was found that the

precipitation of carbides resulted in a drastic

reduction in the hot workability of Nimonic 115. The effect is

illustrated in Fig. 6, which shows the failure during the first

reduction of the chromium carbide containing alloy, while the

alloy with the standard was rolled to the final 3/4" bar. The

carbide morphologies are shown in Fig. 7. The spheroidized MC

carbides aid workability, while the chromium carbides at the

grain boundaries embrittle the alloy. Thus, it was concluded

that in Ni-base superalloys the Cr 23 Cg carbides are not

essential for mechanical properties and, indeed, sometimes such

carbides are detrimental.

Beside carbide effects, chromium also affects the properties

of the matrix since, like cobalt, it partitions to that phase.

Similarly, chromium is a relatively weak solid solution

strengthener compared to the refractory elements. The chromium

effect on gamma prime fraction is also expected to be small,

like cobalt's effect, since the Ni-Al-Cr and the Ni-Al-Co ternary

phase diagrams are quite similar. Certainly a chromium free alloy

would be less prone to sigma formation since chromium is a major

component of that phase. The effects of chromium on stacking fault

P26-6
7

energy and gamma prime anti-phase boundary energy should

certainly be considered, but no intrinsic mechanical problems

would be expected if chromium were removed from a superalloy as

was the case of 903, a chromium free version of alloy 901.

The main reason for the presence of chromium in superalloys

is, of course, the marked improvement in oxidation/hot corrosion

resistance it imparts to the alloys. For hot corrosion resistance

8

an alloy should simply have as much chromium as possible (Fig. 8).

This amount is limited, however, by the 20-25% chromium

6 9 solubility in the matrix '

, beyond which sigma will form.

Since the solubility of chromium in gamma prime is small (about

3-4%), high gamma prime volume fraction alloys are limited in their

total chromium content. Thus, such alloys have intrinsically poor

hot corrosion resistance.

In applications, the superalloys are usually divided into

two groups according to their operating temperatures: the

intermediate temperature range (above 800°C) for turbine

blades. At the lower temperatures fine grain, wrought alloys with

up to about 40% gamma prime are used. At high temperatures, the

strength requirement precludes all but the high gamma prime

fraction, high strength cast alloys. Because of the large gamma

prime fraction, high temperature alloys have less chromium and

must be coated for adequate corrosion protection.

The mechanism of corrosion protection in the complex

P26-7
8

superalloys is not simple. In the case of iron-nickel and cobalt

base alloys forms as the predominant oxide. If the

alloy contains sufficient chromium (Fig. 9), this

layer is continuous and it has a low enough permeability for

oxygen to prevent oxidation of the other metallic species. Such

an oxide is protective by itself. However, if an alloy does not

contain enough chromium (see Fig. 10) , a continuous layer does

not form and other oxides occur instead, such as volatile CoO or

TiC>2 or porous NiO or Fe

2

C> 2

, which are permeable to oxygen.

This oxide situation allows continuous oxidation of the exposed

* 10 surface.

In the case of nickel-base superalloys, the thermodynamically

preferred oxide is A^O^, which has the lowest permeability

of any of the oxides. However, the aluminum content in these

alloys is too low to form a continuous, protective A^O-j

layer. As the oxide starts to form aluminum is drawn to the

surface, dissolving gamma prime and forming a denuded zone marked

'Y' in Fig. 11. Due to the depletion of aluminum in the denuded

zone, other oxides begin to form as fast as oxygen diffuses

through the oxide.

In the nickel-base superalloys, the chromium content is high

enough that also forms in the initial oxidation

reactions. This oxide essentially behaves like a diffusion

barrier for the oxygen and the stable configuration in Fig. 12

is reached. The oxygen activity below the Cr 2°3 layer is

P26-8
9

sufficiently low so that A^O^ alone forms and A1 diffusion

across the denuded zone is the controlling step. This oxide

combination is the source of the excellent oxidation resistance

of the superalloys. Even though an aluminum oxide layer would

provide the best oxidation protection, superalloys require

chromium to form a continuous layer.'*'

0

The layer is irreplaceable for resistance to hot

corrosion in the presence of molten salts. Without this layer

(due to a low chromium content) , the liquid salts and oxygen

preferentially attack grain boundaries (Fig. 13) causing

premature intergranular failure (Fig. 14). The rupture life of

this specimen was considerably shortened by the presence of

NaS0

4 (see Fig. 15)

For this reason low chromium alloys are protected by coating

with corrosion resistant alloys (NiCrAlY, beta, etc.) or plasma

spraying (chromizing or aluminizing) . The compositional

differences between the coating and substrate (such as a beta

phase alloy shown in Fig. 16) form a diffusion couple that can

result in a multitude of intermetallic phases at the interface.

Differences in mechanical and expansion characteristics between

these two alloys are unavoidable. Accordingly thermal fatigue

failure at the interface are likely to occur so the substrate

always contains a considerable amount of chromium as a safety

net.

P26-9
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In summary, chromium can be removed from superalloys without

a severe effect on mechanical properties, but it is essential at

the surface for oxidation resistance especially in the presence

of salt contaminants.

II. Possible Chromium Conservation Options in Superalloys

The current solution for resisting hot corrosion in

superalloy design is the addition of as much chromium as the

matrix can hold without forming sigma phase and then coating the

alloy if that amount is not sufficient (see Fig. 17). Aside from

the conservation consideration that over 90% of the chromium is

not needed, the presence of this chromium, in fact, can

deleteriously affect the alloys. As mentioned, these negative

aspects include sigma and mu phase instabilities, excessive

m

23C

6 Precipi- tat i°n r and relatively high thermal expansion

coefficients.

The chromium requirement of superalloys can be redefined by

three criteria:

1. The near surface of the alloy (a minimum of 10 microns)

must contain sufficient chromium for corrosion resistance (~'20%).

2. This layer must be integrally bonded to the substrate

with no continuous intermetallic layers or physical interfaces.

3. The configuration must be sufficiently stable over prolonged periods ( ^ 1000 hours). We performed a simulation of the

P26-10
11

diffusion process (Fig. 18) and determined that the initial 10

micron profile would be stable at 800°C, but a deeper profile

of 50-100 microns would be required for 1000 hours of operation

at 1000°C (Fig. 19).

The available techniques for producing surface chemical

changes are:

1. Claddings and coatings—Using current diffusion technology, chromium-free base alloys can be coated or clad by

conventional techniques. However, this approach has been

rejected, to date, since a crack forming in the coating would

result in disastrous substrate degradation in hot corrosion

environments. Currently all base alloys that are coated contain

at least 8% chromium for insurance.

2. Plasma spraying and overlays—Overlays also have a distinct interface and the same failure consequences as claddings

or coatings.

12

3. Dual composition powders —By using two different powder

compositions for the interior and surface of the disk components,

a sizable amount of the chromium could theoretically be saved.

However, to our knowledge, this technique has still not been

successfully used to produce a reliable crack-free component.

13 14

4. Pack chromizing ' —This technique has been successfully used to reproduce a surface alloy on stainless steels.

The steel component, for example, is usually packed in a mixture

ferro-chrome or chromium metal and alumina powders with a small

amount of ammonium chloride and annealed in an inert or reducing

P26-11
12

environment until the required diffusion profile is achieved.

The gaseous chromium chloride is the carrier exchange gas of the

chromium atoms rather than the metal/metal interface as is the

case for the true diffusion couple. Currently this technique has

been used experimentally to apply coatings to superalloys, but

most of the research is unpublished company proprietary or

published in the Soviet Union. This technique might be applied to

achieve a smooth concentration profile on the surface of a chromium-free alloy by adjusting chromium and nickel activities

in the pack or the exchange gas. However, we do not know of any

success, to date, of producing a significant surface concentration of chromium without forming an alpha chromium layer.

5. Laser glazing^'^ or electron beam surface welding

Laser glazing is a rapid solidification technique that

has been used to construct disk shapes from layers of nickelmolybdenum alloys roughly 100 microns thick. Conceivably this

technique could be used to apply a "layer of a chromium

containing alloy to a chromium-free superalloy substrate.

However, as yet, no conventional alloy has been laser glazed

without forming surface cracks.

17 18 19

6. Chromium ion implantation ' ' —Ion implantation is

readily used as a surface alloying technique for low temperature

wear or for aqueous corrosion resistance. However, the depth of

the alloy layer of chromium is on the order of .1 micron, which

is well below the 10 micron minimum required. Additional difficulties include sputtering loss and a low beam current (measured in

microamps). Hence, straight ion implantation can be eliminated

P26-12
13

at this time as a viable alternative.

19

7. Ion recoil implantation —As an alternative to using

chromium ions as the source, recoil implantation can conceivably

use a vapor or electrolytically deposited layer (of chromium)

as the target for a high current (milliamps) beam of argon ions.

The atom localized impact can in theory drive the chromium into

the surface to a depth of approximately 1 micron. Multiple

anneals and reapplication (Fig. 2o) could result in a substantial

profile to a depth of 10 microns or more with no sharp chromium

alloy interface.

Isothermal annealing of the implanted layer would most likely

produce a smooth concentration profile following Fick's law. Since

no diffusion barrier exists at the 10 micron depth, repeat implantation/anneal cycles will distribute previous implant layers over

an ever-increasing range before a substantial concentration is

developed at the 10 micron depth. However, we propose that the

diffusion range can be restricted by the application of a steep

thermal gradient. Controlled laser annealing20 ' 21 is such a technique. Laser energy can be concentrated to within a few

microns of the surface and thus constrain chromium diffusion in

the near surface layer. With this method a high concentration

near-surface layer can be built with a minimum number of implantation cycles.

The same recoil mechanism may be used to bond a thicker

chromium layer to the alloy surface. The argon ion beam disrupts

P26-13
14

the interface yielding a smooth transition layer for the

diffusion couple without the problem of decohesion. In this

alternative scenario, laser anneal can again serve as the

spatially contained heat source for driving the chromium rapidly

through the interface resulting in a chromium bearing surface

alloy.

In summary, we assert that the chromium metal in superalloys

is a critical strategic material which is quite vulnerable to supply side fluctuations. We have shown that chromium is essential

for the elevated temperature oxidation and hot corrosion resistance of these alloys. However, the high chromium content

required for the surface protection is not needed for the

mechanical properties of the superalloys and, indeed, is often

considered more detrimental than beneficial for these properties.

Accordingly, we have reviewed the various methods of achieving the

srface requirement of chromium while eliminating the bulk of the

chromium from the alloys. While current methods of protecting low

chromium alloys have drawbacks when applied to chromium-free

alloys, several alternate methods of surface alloying, including

ion implantation and laser annealing, have recently shown possible applications in the area of conservation of chromium in

superalloys.

P26-14
15

hi. &efeL£nggg

1. J.L. Morning, N.A. Matthews and E.C. Petersen, in Mineral Facts

and Problems 1980 , D.S. Bureau of Mines Bulletin 671, 1980,

pp. 167-182.

2. U.S. Bureau of Mines Commodity Data Summary, "Chromium," 1980-1981.

3. L.R. Curwick, W. A. Petersen and H.V. Makar, U.S. Bureau of Mines

IC #8821, 1980.

4. Private communication, W.J. Boesch, Special Metals Corporation,

New Hartford, New York.

5. Private communication, A. FitzGibbon, Elkem Metals, Pittsburgh.

6. R.N. Jarrett and J.K. Tien, Met. Trans. A, Vol. 13a, June

1982, pp. 1021-1032.

7. J.K. Tien and R.N. Jarrett, in High Temperature alloys for

Gas Turbines 1982, eds. R. Brunetaud et al., D. Reidel Publishing

Co., Dordrecht, The Netherlands, pp. 423-446.

8. The Superalloys . eds. C.T. Sims and W.C. Hagel, John Wiley and

Sons, Inc., New York, 1972.

9. O.H. Kriege and J.M. Baris, Trans. ASM, Vol. 62, 1969, pp. 195-200.

10. F.S. Pettit and J.K. Tien, in Corrosion Fatigue . 1972, pp. 576-

589.

11. J.K. Tien and J.M. Davidson, Advances in Corrosion Science and

Technology, Vol. 7, 1980, p. 1-51.

12. Private communication, J. Stulga, Crucible Inc. Research Center,

Pittsburgh.

13. R. Sivakumar, Trans. Indian Inst, of Metals, Vol. 33, No. 5, 1980,

pp. 398-403.

14. R. Sivakumar and L.L. Seigle, Met. Trans. A, Vol. 7A, 1976, p.

1073.

15. E.M. Breinan, B.H. Rear, C.M. Banas and L.E. Greenwald, in

Superalloys; Metallurgy and Ma nufacture . Claitor’s Publishing

Div. , Baton Rouge, 1976, pp. 435-450.

16. D.B. Snow, E.M. Breinan and B.H. Rear, in Superalloys 1980 . ASM,

Metals Park, 1980, pp. 189-203.

17. A.B. Campbell III, B.D. Sartwell and P.B. Needham, Jr., U.S. bureau

of Mines Report RI #8387, 1979.

P26-15
16

18. C.W. Draper, Journal of Metals, June 1982, pp. 24-32.

19. Private Communication, D. Moon and W. Nahemow, Westinghouse Research

and Development, Pittsburgh.

20. G. Dearnaley, Journal of Metals, Sept. 1982, pp. 18-28.

21. Private communication, S. Copley, University of Southern California,

Los Angeles.

P26-16
17

Other

Fig. 1 1980 Total U.S. chromium consumption

(849 million pounds)

Nuclear

Fig. 2 Superalloy use of refined chromium metal

P26-17
RELATIVE

AMOUNT

18

Fig. 3 Microstructure of a typical nickel-base superalloy

with carbides at the grain boundaries

1.0

0.8

0.6

0.4

0.2

0 .

0

0 5 10 15

WEIGHT X COBALT

Fig. 4 Effect of cobalt on carbides in Udimet 700

P26-18
STRESS

(in

MPa)

19

Fig. 5 Room temperature tensile properties of cobalt

modified Udimet 700

P26-19
20

Nimon\cll5

OY. CobalV

1 2 3 l S l> '7; H 1. W W V2

nt •< vie o* vie «J« «b <>•* alt vji J* •’* A » ^

i «» > » t , v.t, * . * * — •

Fig. 6 Failure of low cobalt Nimonic 115 during rolling

due to chromium carbide content

Fig. 7 Carbide microstructure of standard N115 (right) and

low cobalt N115 (left)Note the nearly continuous

chromium carbides in the cobalt-free alloy.

P26-20
21

1 IN- 1 00

\

\• U-700

\

\» U- 500

y RENE-41 • HASTELLOY X

_ M3 1

\HM-30 N, ' MCr

• J I I I I 10 20 30 40 30 WEIGHT %Cf

Idealization of the effect of chromium on the hot-corrosion resistance of

nickel-base alloys in burner-rig tests.

ALLOY Cr Content y' fraction Cr ( in y) Cr ( in y

# ) Wrought

Waspaloy 19.5 20. 23. 2.

Udimet 700 15.0 45. 22. 3.

Nimonic 115 15.0 50. 24. 4.

Cast

Inconel 713c 12.5 50. 23. 3.

IN 100 10.0 65. 22. 3.

MarM 200 9.0 55. 17. 3.

Fig. 8 Phase compositions (9 , 6 ) and the effect of chromium

on hot corrosion resistance (8) . Note the chromium

solubility in the the gamma and gamma prime phases

of all the alloys as nearly constant. This limit

is determined by the sigma phase boundary in the

multicomponent phase diagram.

P26-21
22

Fig. 9 iron-chromium alloy with protective

chromium oxide layer

Fig. 10 Oxidation of a low chromium alloy

P26-22
23

Fig. 11 Formation of a gamma prime denuded zone below

the alumium oxide in a nickel-base superalloy

SUPERALLOY OXIDE SCALE

Fig. 12 Superalloy oxide protective scales

P26-23
24

Fig. 13 Preferential grain boundary attack in the hot

corrosion of a nickel-base superalloy

P26-24
STRAIN

°/<

25

Fig. 14 Intergranular failure of a hot corrosion/

stress rupture specimen

TIME (HOURS)

Fig. 15 Corrosion induced premature failure of the

above specimen

P26-25
26

ZONE 1 0(NiAI) matrix bcc a Cr,

Mo substrata phases

IQhLA

single phase (KNiAl), Cr, Mo,

Ti, Co in solution ZONE 3 3(NiAf) matrix, TIC, M C , and o(Cr,Mo,Co) 236

o phase In coating affected substrate

Fig. 16 Diffusion boundaries of a beta phase

coated superalloy

P26-26
27

CONVENTIONAL SOLUTIONS

MCrAlY

T +T*

15 % Cr

T+r*

10 X Cr

CONSERVATION OPTIONS

Cr-rich Alloy MCrAlY

Cr-rich Allov

0% Cr 0% Cr

10 urn Implanted

With 15% Cr

MCrAlY

Cr Implanted Layer ,

0% Cr 0% Cr

Fig. 17 Methods of providing superalloys with

adequate corrosion resistance

P26-27
28

c(x,t)

2(2tt)*S(1+t) 3/2 “p[ 2nrn]ERF C[ }2(x(i+T) i] + exp[ ~2;ht) ]erf V2(t(1+t)

^

WHERE:

2Dt

52

a = xp

6

$ = total ion dose

.0

AND THE INITIAL PROFILE IS:

c(x,o) exp[- ( x~ xo ) 2

1 2«2

J

Fig. 18 Diffusion profile of a surface alloy

in a nickel-base alloy

P26-28
CHROMIUM

CONTENT

CHROMIUM

CONTENT

29

30,_

DIFFUSION PROFILES

Initial Depth lOum

Exposure for 1000 hrs

4 8 12 DEPTH (in um)

10

Initial Depth 50um

1000 hrs at 1000 C

OL j * L_ 4 8

DEPTH

_J I L 12 (in um)_L

16 20

Fig. 19 Required profiles for 800°C and 1000 C exposure

P26-29
CHROMIUM

CONTENT

CHROMIUM

CONTENT

CHROMIUM

CONTENT

CHROMIUM

CONTENT

30

Specimen

Surface

Fig. 2o Multiple implantation approach

P26-30
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.

DEVELOPMENT OF MODIFIED 9 Cr-1 Mo STEEL1

Vinod K. Sikka

Metals and Ceramics Division

Oak Ridge National Laboratory

Oak Ridge, TN 37830

The status of development and commercialization of a modified 9 Cr-1 Mo alloy is presented. The alloy is modified

by the addition of 0.06 to 0.10% Nb and 0.18 to 0.25% V. The

alloy is recommended for use in the normalized and tempered

condition (1038°C for 1 h, air cooled to room temperature;

760°C for 1 h, air cooled to room temperature). Heat treatment, Charpy impact, tensile, and creep properties of the

alloy are described in detail along with brief description of

other properties.

The modified alloy has creep strength that exceeds that

of standard 9 Cr-1 Mo and 2 1/ A Cr-1 Mo steels for the temperature range from 427 to 704°C. The total elongation and

reduction of area values for all test temperatures and rupture

times up to 22,500 h exceed 15 and 70%, respectively. The

estimated design allowable stresses for this alloy are higher

than those for standard 9 Cr-1 Mo and 21/4 Cr-1 Mo steel. At

550° C and above, these values are twice those of the other

alloys. Operating experience on this alloy is being obtained

by installing tubes in various steam power plants.

1. Introduction

A 9 Cr-1 Mo steel with properties improved over the 21/4 Cr-1 Mo

steel and that of other ferritics in the ranges from 9 to 12% Cr and

1 to 2% Mo has been developed recently [1, 2]. The development of this

alloy was funded jointly by the U.S. Department of Energy fossil energy

Research sponsored by the Office of Fossil Energy and the Office of

Breeder Reactor Technology Project, U.S. Department of Energy, under

contract W-7405-eng-26 with Union Carbide Corporation.

P27-1
and breeder reactor development programs. The purpose of this paper is

to describe the status of this development, with special emphasis on the

status of approval of this material in the ASTM Specifications and the

ASME Boiler and Pressure Vessel Code .

2. Chemical Specifications

The composition specifications of modified 9 Cr-1 Mo steel are listed in Table 1 and compared with those of standard 9 Cr-1 Mo steel.

The main difference with the modified alloy compared with the standard

alloy include

1. addition of niobium and vanadium,

2. a specified range for each element, and

3. a specification for nitrogen, which is not listed for the standard

9 Cr-1 Mo.

Table 1. Chemical analysis of modified

9 Cr-1 Mo steel and its comparison

with standard 9 Cr-1 Mo steel

Content range, wt %

Element Modified Standard

9 Cr-1 Mo 9 Cr-1 Mo

(Grade 91) (Grade 9)

Carbon

Manganese

Phosphorus

Sulfur

Silicon

Chromium

Molybdenum

Nickel

Vanadium

Niobium

Nitrogen

Aluminum

0.08-0. 12

0.30-0.60

0. 020 max

0. 010 max

0.20-0.50

8.00-9.50

0.85-1.05

0. 40 max

0. 18-0.25

0.06-0.10

0.030-0.070

0. 04 max

0. 15 max

0.30-0.60

0. 030 max

0. 030 max

1.00 max

8.00-10.00

0. 90-1. 10

P27-2
Both niobium and vanadium are added to the alloy to improve its

elevated-temperature strength properties. Microstructural work has indicated that the improved strength of the modified alloy comes from two

factors. First, fine M2 3 C6 precipitate particles nucleate on Nb(C,N),

which comes out first during the heat treatment. Second, the vanadium

enters M23 C0 and retards its growth at the service temperature. The

finer distribution of M23 C5 adds to strength, and its retarded growth

holds the strength for long periods of time at the service temperature.

3. Heat Treatment

The alloy is recommended for use in the normalized and tempered

condition. The normalizing treatment consists of heating the alloy to

1040°C, holding for 1 h for material up to 25 mm thick, and then air

cooling to room temperature. This treatment produces a fully martensitic

structure. The typical hardness in this condition is Rockwell C40. The

tempering treatment consists of heating the normalized steel to 760°C,

holding for 1 h up to 25-mm thickness, and then air cooling to room

temperature. The typical hardness in this condition is Rockwell B95.

Optical and transmission electron micrographs of tempered martensite are shown in Figs. 1 and 2. Figure 1 shows that the alloy is single phase

Figure 1. Typical microstructure of modified 9 Cr-1 Mo steel after the

nominal normalizing and tempering treatment (1040°C for 1 h,

760°C for 1 h). Note that the grain size is very fine.

P27-3
Figure 2. Transmission electron micrograph of a specimen normalized ~at

1040°C for 1 h and tempered at 760°C for 1 h. The micrograph

shows that tempered martensite consists of dislocation

substructure and carbides (M23 C0 and MC) on both the grain

and subgrain boundaries.

(free from 6 -ferrite) and has a fine grain size (ASTM 8—9). The transmission electron micrograph shows that the tempered microstructure has

high-dislocation-density subboundaries in the matrix. The subboundaries

are stabilized by the precipitation of carbides on them. Carbides also

precipitate at the prior austenite grain boundaries.

The tempering response of the alloy can be described by the

Hollomon-Jaf f e (HJ) tempering parameter. This parameter is given by

HJ = T(]C + log t) , ( 1)

P27-4
where

T = temperature in kelvins,

t = time in hours,

C = HJ constant.

The optimized value of C for two commercial heats (30176 and 30394)

characterized to date was determined to be 22.3.

The correlation of tempering parameter with hardness, Charpy

V-notch energy, 0.2% yield strength, and total elongation at room temperature is shown in Fig. 3. The 0.2% yield and ultimate tensile

strengths are well correlated with room-temperature hardness (Fig. 4).

Such correlations are useful when the material has to meet both the

tensile property and hardness criteria.

The grain coarsening response of the alloy was examined as a function of the normalizing temperature and time (Fig. 5). Grain coarsening

of the modified alloy occurred only when the normalizing temperature was increased about 100°C or more above the specified normalizing temperature of 1040°C. Even in the coarsened condition, the grain size for the

modified alloy remained near ASTM No. 5.

4. Commercial Melting and Fabrication

Eight heats of this alloy have been melted by commercial vendors

(Table 2). Many vendors have fabricated this alloy (Table 3). The

overall consensus is that melting and fabrication of this alloy presents

no technical problems, although costs will obviously vary depending upon

the melting and fabrication process selected.

5. Mechanical Properties

The main emphasis on the mechanical properties of modified alloy

was limited to Charpy-impact , tensile, and creep testing. Each of these

properties is described here briefly.

5. 1 Charpy-impact Properties

The curves of Charpy V-notch impact energy versus test temperature

for two commercial heats of modified 9 Cr-1 Mo steel are compared with a heat of standard 9 Cr-1 Mo steel in Fig. 6. All tests were conducted on

P27-5
T

— c*x sn

M O —* a 2 *

2 £

“ 2

t A*'

s

AN

S

s —

O w O CC

3 ^ ” J J5 K * • <£111111

- O OO D0O0

2 i

“ s V

- yi itj-

% .#*

J L

(i ioi czz) i - m

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•u T33

cO

.3u

603

<U3U

CO •a

i—

(U •H

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• o

5*603

a>3

cuCo

a.3

CO X!O

CO

CO

<03

T3

(JcO X

m

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3o

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CO 0

cO <0

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<u o

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cO3 4-1 cO cO a

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3 *H

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a) i—

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CO

<u3

3

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Cu

P27-6
HARDNESS (ROCKWELL C) 6 8 10 12 14 16 18 20 22

HARDNESS (ROCKWELL B)

o

Q.5

X

h-o

2

UJ

a:

>

(VJ 6

(b)

HARDNESS (ROCKWELL C)

6 8 10 12 14 16 18 20 22

HARDNESS (ROCKWELL 8)

oa

5

o

2

UJ

crhtoUJ

UJK

<

5

3

Figure 4. Room-temperature strength versus hardness for two commercial

heats of modified 9 Cr-1 Mo steel. (a) Yield. (b) Ultimate.

P27-7
Hj « T 122.3 log t ) Figure 5. Grain coarsening behavior of modified 9 Cr-1 Mo steel.

Table 2. List of commercial heats melted on modified 9 Cr-1 Mo steel

Heat Melter Heat size

(tons) Melting practicea

F5349 Quaker 4 AOD

30383 CarTech 15 AOD

3039 4^ CarTech 15 AOD/ ESR

30176c CarTech 15 AOD/ESR

30182c CarTech 15 AOD/ESR

10148^ Electralloy 15 AOD and AOD/ESR

XA3602 Combustion

Engineering

0.5 Air induction

91887 CarTech 2.5 Electric/ESR

YYC982C Sumitomo, Japan 2.0 Vacuum induction

59020 NKK, Japan 5.0 Vacuum induction

aAOD = argon-oxygen decarburization. ESR = electroslag remelting.

^Heat 30394 is half of heat 30383 and was electroslag remelted.

^Heats 30176 and 30182 are the same heat. Their numbers are different because CarTech identifies ingots separately.

^Heat 10148 was melted by AOD process. Half this heat was subsequently electroslag remelted.

P27-8
Table

3.

Summary

of

information

on

fabrication

of

commercial

heats

of

modified

9

Cr-1

Mo

alloy

cd 1 <u O >< x x X X

•H a o O. o e o c a <u U A o 0) O a) <u <u tu o X

X co H X CO H X x H X M CO X X X

c0 Z CO P td CO P td s p td 0 CO z 5s x

Pm X a) td 0 a; nj a •H X td 0 •H CD Ot! tu o u < u < H U o -3 H *“) o o PQ

•X3 x X X X X ”3 <U tu tu <u tu <u — O rH iH 1—

t

i—i rH 1—I X

X iH r-t 1—

1

rH t—1 i—4 <u

4-1 o O o O X o X o p

<u P P p P tu p <u p tu S u u CL

4-1 4-J 4-1 4-1 p 4J p 4-1 0 o O O o o <u o <u O tu o X X X X •H X •H X 4-J x

•H X CL X CL <u X 4J A A * <u 0k <u — <u 0k — i-H 0)

td T3 T) T3 X X X X >> X X X X X X >H X X

o ai <U tu 3 tu <U 3 P tu <u 3 P <u <U O 3 <u •H t>0 60 60 P 60 60 P td 60 60 P td 60 60 P p i—4 P p Li P 4-1 P P 4-J 4J P P 4-J 4-J P P 4-J iH X o O O X o O X o o o X o o O X X o

td t4H 4-1 44 (U 44 44 <u p 44 44 ai p 44 44 >H tu p X O

4-1 4-J 4-J 4-> 4J 4-> 4-1 4-1 4-1 4J 4-> 4-J 4-J 4-J O 4-> 4-J O o O o o O o o O O o o o o o o X X X X X X X X X X X X X X X X

vo »-h m in

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60 CD X X X

3 O X co c/d •H •H x X x

X 4-J 1 1 1 iH o Q Q Q o Q Q o

<U cO q O O o O q o

S V4 <! <3 <3 <3 c <3 <3 a

X X X X X X

co X X X X X

X1 X1 X1 X1 X1 X1 Q

1 Q Q Q Q Q Q Q Q O O o O O O O O O

<3 <3 <3 <3 <3 c <3 C <3

P

tu >

X o o o X f—

t

rH rH

tu rH rH rH X X X X X X X X X X X X X td CO CO p o u u o o CJ o u u u u o P U P

<u <u tu (U tu <u tu <u tu (U <u <u <u X 4-J 4-J X H H H H H H H H H H H H o o a

td P P P P P P P P P P P P tu CD CD 5- td td td td td td td td td td td td i—t rH rH u O o o c_> c_> o V V u o (_> X P3 W

4-J c^ X CO CO <r CM CM CM CM v£> 00 oo 00

cO <t 22 00 O' O' On ON 00 00 00 00 r>. <r -o -o

CD co X PO CO CO X co X r—H rH rH rH r—M rH rH rH

PQ in o o o o o o o o o o o o o o o

Pm X CO CO co X CO X CO CO X X CO rH r"H rH

P27-9

10148

Electralloy

AOD

Pipe

245

25

Hot

pilgered

Phoenix
Table

3.

(Continued)

CO

u CO 3 a; u rH CO rH CO rH co »H co rH

03 03 CM CO CM (0 CM 03 CM c o o J CO o C0 o co o CO o •rH B B

•H >-i rH 2 rH P rH rH ctf r-i o O U CD o <0 CJ <0 cj 0) u 4-» 2 2 2 > >1 > >i > >1 > >i CO 2 •rl •H

03 •H cj iH CJ iH cj

u N—

B s txS as c s c c M w 3 3 JkS 2 2

2> 2 ZD

H o cn C/3 Z z z 3 3 3 3 3 3 3 3 X cO 3 2 cO 3 X C Mi 2 a> 2 2 3 03 2 2 rH 2 2 2 rH 2 X 4-4 2 2 o 2 2

a) rH CO rH rH

U rH rH s O 03 o O O o o a 2 CJ o 4-J o u

ai o 5 0k Mi 0k — 0k

—

•H 2 rH 01 X) 2 X 2 U <U rH 60 <D 4) * CD 01

03X 2 X)X 2 03 rH

X 2 XX 2 j CD CO 01 (1) 3 2 60 iH 3 CD 3 3

iH rH rH 60 00 Mi Q1 3 2 >H 2 00

u 2

1-1 rH i-H Mi 4-1 CJ MH 4-J 2 >Hu 2 2 o O O O X 3 •H 2 X X o X X

tO >H IM UM MH ai 2 *H rH <1> a> 14H CD ai b Q1 4-J o

4-> 4-1 4-1 4J 4-1 Mi C2

O 4-J 2 4J 4-) 2 o O o o o a; O O o O o

32 2 2 32 2 o 32 2 32 2 - CO § CO

<0 CO CO e <r o -3- -O' 0) • • • • N o O' 00 CO vO m co vD

•H •H rH CNI rH

CO 2H

4-J O3rr—

o CO rH o C\1 <? rH 2 H 2 rH

1-1

Q o LO o co in

m r-- UH I-- U1

CmO CNI CN CM

u s 1 J o 3 a; 01 60

XX a) X) 4-) 4-1 CO o C <l) ai HI 0) 4-J 01 01 D 03 cO 4-1 2 3 2 2 2 2 03 2 2

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iH 1-1

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<0 cO

O o O

OO •H

M MH HHM P 2 u c <c < <5 < < > > > > >

CM

CO CO CO CO CO a CM CM CM 2 o o o o O

i-H 1—rH rH rH

U

u CJ

a CJ Sh >1 >N >>

cj

a CJ CJ cj 60 4 — "—

~—

3

01 >3 rH >1 rH rH >1 rH >3 rH iH

4-1 o cfl

O cfl o CO o cO o c« e Mi f—

rH CO rH CO rH CO T—CO rH CO o 0)

CD rH !-> rH 2 rH Ml rH Mi rH Ml iH 01O o 2 cO 01 CO CO 03 01 03 ai 03 0) 2 C B B

1-1 > IM > P > u > U > CO •H o o

4-1 •H 4-1 -rH 4J •H 4J •H

u iH 3 60 2 2

CJ 3 o s O c a c O c 2 3 •H iH

ai 2 (0

p CD 2 a; 2 CD 2 B

w

B S as i—

rH rH rH rH o 3 3 as

X

w

w

W

w

wu C/3 C/3 z 2

CJ CJ

CNI CNI CNI 4J 00 oo 00 00 00 o 2 00

O o o

03 -3- <r <T St 2 O' O' CN CN CN

a) rH rH rH rH rH ro 2 CJ

OO

O

X o o o o o <c >• >4 O' On

rH rH >4 f-1 m 1T|

m

P27-10

a

AOD=

argon-oxygen

decarburization,

ESR=

electroslag

remelting,

VIM=

vacuum

induction

melting.
200

TEST TEMPERATURE ( °C ) Figure 6. Effect of melting process on the Charpy V-notch toughness

from two 14-Mg (15-ton) commercial ESR heats of modified

9 Cr-1 Mo steel and one standard 9 Cr-1 Mo steel plate.

specimens with the V-notch parallel to the rolling direction. The

standard 9 Cr-1 Mo heat was electric furnace melted. One modified

9 Cr-1 Mo steel heat was melted by both the AOD (argon-oxygen decarburization) and AOD-ESR combination. All the second heat of modified

9 Cr-1 Mo was melted by the AOD-ESR process. Data in this figure show

that the 68 J (50 ft-lb) ductile-to-brittle transition temperatures for

both heats of modified 9 Cr-1 Mo are lower than that for the standard

9 Cr-1 Mo heat. The upper-shelf energies for the AOD-ESR-melted

modified 9 Cr-1 Mo heats are much higher than those for either

AOD-melted modified 9 Cr-1 Mo steel or electric-furnace-melted standard

9 Cr-1 Mo steel. This figure also shows that Charpy properties are affected by both the silicon content and the melting practice [3]. The

lower silicon content or electroslag remelting of the material produces

lower transition temperatures and higher upper-shelf energies. However,

between the two factors, melting practice appears to have a much more

potent effect than silicon content.

P27-11
5.2 Tensile Properties

Tensile properties of commercial heats of modified 9 Cr-1 Mo steel

have been measured for test temperatures in the range from room temperature to 760°C. Data include the effects of melting practice,

compositional differences, strain rate (8.0 to 0. 00008/min) , postweld

heat treatment (1—112 h at 732°C), isothermal annealing treatment

(1038°C for 1 h, fast cool to 704°C, and hold for 14 h, followed by

air cooling), and thermal aging treatment (5000 and 11,600 h at 538,

593, and 649°C).

Typical plots of 0.2% yield and ultimate tensile strengths against

test temperature are shown in Fig. 7. Similar plots for total elongation and reduction of area are shown in Fig. 8. These plots combine

data on plate, tube, and bar. All tests were at a nominal strain rate

of 0.004/min. Room-temperature data on individual product forms were

used to determine the average values, the standard deviation in the

room-temperature data, and the recommended minimum properties for use in

ASTM specifications. These values are listed in Table 4, and the recom- mended minimum values are the same for all three products. The best fit

average and minimum curves based on the room-temperature specified

minimum properties (see Table 4) are also included in Figs. 7 and 8.

These plots show that the minimum specified properties at room temperature are conservative enough to include data on all commercial heats

through test temperatures of 760°C.

Tensile data anlysis also showed the following features.

1. Melting practice, AOD versus AOD-ESR, had no effect on tensile

properties for test temperatures up to 760°C. This is different from

impact properties, which were far superior for AOD-ESR-melted material

to those for AOD-melted material.

2. The tensile properties of modified 9 Cr-1 Mo steel are sensitive to strain rate at room temperature and three higher test temperatures: 427, 538, and 649°C. The yield and ultimate tensile strengths

decrease some with decreasing strain rate at room temperature. This

P27-12
(b) Temperature (°C)

Figure 7. Plots of (a) 0.2% offset yield strength and (b) ultimate

tensile strength as functions of test temperature for plate,

bar, and tube of commercial heats of modified 9 Cr-1 Mo

steel. The best fit average and minimum curves based on room-temperature specified minimum yield and ultimate tensile

strength of 414 and 586 MPa are included also in these

figures.

P27-13
o

Figure 8. Plots of (a) total elongation and (b) reduction of area as functions of test temperature for plate, bar, and tube of

commercial heats of modified 9 Cr-1 Mo steel. The best fit

average and minimum curves based on the room-temperature

specified minimum total elongation and reduction of area of

18 and 55% are included also in these figures.

P27-14
Table 4. Minimum tensile properties at room temperature specified

for various products of modified 9 Cr-1 Mo steel tested

in the normalized and tempered conditiona

Product Calculation

Strength, MPa Total

elongation

(%)

Reduction

of area

0.2% Ultimate (%>

Plate Average 545 683 24.7 69.8

Average — 2SD 496 634 20.6 61.6

Average — 3SD 476 607 18.5 57.5

Recommended min 414 586 18 55

Tube Average 524 669 27 74

Average — 2SD 483 627 23 71

Average — 3SD 462 607 21 70

Recommended min 414 586 18 55

Bar Average 531 676 28 73

Average — 2SD 469 614 26 70

Average — 3SD 434 586 25.5 68

Recommended min 414 586 18 55

aAt 1038°C for 1 h, air cooled to room temperature, and at 760°

C

for 1 h, air cooled to room temperature.

occurs without any change in ductility. At 427 and 538°C strength

decreases and ductility increases with decreasing strain rate. At

649°C, strength decreases with some decrease in ductility.

3. A decrease of 14 to 28°C in tempering temperature from 760°C

substantially increased yield and ultimate tensile strengths for test

temperatures through 427°C. At a test temperature of 649°C the increase

in strength was relatively small. Ductility was only slightly affected

by this drop in tempering temperature. Changes in hardness and Charpyimpact energy as a result of change in tempering temperature are shown

in Fig. 3.

4. Tensile tests were conducted at room temperature and at the

aging temperature on a single heat (F5349) aged for 5000 and 11,600 h at

538, 593, and 649°C. Only small changes were observed in properties as

a result of this treatment.

P27-15
5.3 Creep Data

A substantial number of creep tests on the commercial heats of

modified 9 Cr-1 Mo steel are either ruptured or in progress. These

tests are being run at 427, 454, 482, 538, 593, 649, 677, and 704°C.

Creep rupture data at 593°C on commercial heats of modified 9 Cr-1 Mo

steel are plotted in Fig. 9. Data in this figure are compared with the

average curves for the modified 9 Cr-1 Mo, standard British 9 Cr-1 Mo,

standard American 9 Cr-1 Mo and 2 1/4 Cr-1 Mo. The modified and

standard British 9 Cr-1 Mo steels are both tested in the normalized

and tempered condition. The standard American 9 Cr-1 Mo and

21/4 Cr-1 Mo steels were tested in the annealed condition. The

following observations are derived from Fig. 9.

1. The modified 9 Cr-1 Mo steel is significantly stronger than

standard 9 Cr-1 Mo and 21/4 Cr-1 Mo. The strength difference becomes even more important for rupture times exceeding 1000 h.

400

200

100 a 80

5

{/) 60

UJa£ 40

20

(0° 10* tO

2 I0

3 I0

4 I0S TIME TO RUPTURE ( h ) Figure 9. Stress-rupture plot at 593°C comparing data on commercial

heats of modified 9 Cr-1 Mo steel (tested in normalized and

tempered condition) with the average curves for the standard

British 9 Cr-1 Mo steel tested in the normalized and tempered

condition and standard American 9 Cr-1 Mo steel and

21/4 Cr-1 Mo steel tested in the annealed condition.

i—rrri—ittt|

i—rrrf—i—

m

-

]—i ttt

-- MOOlPIED 9Cr-1Mo; NT: AVG STANDARD BRITISH 9Cr-IMo; N T ; AVG

2 1/4 Cr - I Mo ; A : AVG STANDARD 9 Cr - 1 Mo.- A; AVG

V -30394 CARTECH AOCVESR O - 10148 ELECTRAULOY AOD

D> - 91887 CARTECH ESR —INDICATES TEST IN PROGRESS

TEST TEMPERATURE 593"C (1100 Nr • NORMALIZED AND TEMPERED

A ANNEALED

I i i i I I I I I I I I I I I I I ] 1 1 I 1—LL

P27-16
2. Changing the heat treatment condition from annealed to

normalized and tempered improves the creep-rupture strength only a small amount. However, for the same heat treatment (normalizing and

tempering) the composition modification made in our alloy produced a very large improvement in creep-rupture strength.

Figure 9 compares rupture strength only at 593°C. The 1

0

5 —h creeprupture strength at a range of temperatures (427—7 32°C) is compared for

various alloys in Fig. 10. This figure also includes the curve for

21/4 Cr-1 Mo steel tested in the normalized and tempered condition.

This figure shows again that the modified 9 Cr-1 Mo alloy is stronger

than standard 9 Cr-1 Mo and 21/4 Cr-1 Mo steel for the entire temperature range. The strength differences become more significant for

temperatures exceeding 500°C.

TEMPERATURE CC)

Figure 10. Variation of lO^-h creep rupture strength with temperature

for several materials.

P27-17
When the creep strength of an alloy is increased, one is concerned

with the associated loss in ductility. The total elongation and reduction of area for creep tests ruptured thus far on the commercial heats

of modified 9 Cr-1 Mo steel are plotted in Fig. 11. The total elongation for most heats exceed 15%, and the reduction of area exceeds 70%.

Furthermore, these data show no indication of drop in ductility with

increasing rupture time. This is true for all tests temperatures.

The data presented above show that the alloy we have developed has

high creep strength and excellent creep ductility.

5.4 Estimated Design Allowable Stresses

Tensile and creep data were employed to estimate the design allowable stresses for the modified 9 Cr-1 Mo steel. The procedure employed

for these estimates was the one used for the ASME Code, Sect. VIII. The

maximum design allowable stresses estimated for the modified 9 Cr-1 Mo

steel (normalized and tempered) are compared with those for 2 1/4 Cr-1 Mo

(annealed and normalized and tempered) and standard 9 Cr-1 Mo steel

(normalized and tempered) in Fig. 12. This figure shows that the design

allowable stresses for the modified alloy are higher than those of

standard 9 Cr-1 Mo and 21/4 Cr-1 Mo steels for the entire temperature

range. Furthermore, at 550°C and above the design allowable stresses

become approximately twice those of the other alloys.

6. Related Properties

Data on both physical and mechanical properties other than these

presented in this paper are either available or being collected in an effort to completely characterize the modified 9 Cr-1 Mo alloy. Each of

these properties is listed with its status described briefly.

6.1 Physical Properties

In this area effort is under way to determine the thermal expansion

coefficient, thermal conductivity, and heat capacity of one heat each

of standard and modified 9 Cr-1 Mo alloy. The objective of the study

is to determine if there are significant differences in the physical

properties of the two alloys.
REDUCTION

OF

AREA

(%)

TOTAL

ELONGATION

(%)

Figure 11. Plots of (top) total elongation and (bottom) reduction of

area as functions of rupture time for commercial heats of

modified 9 Cr-1 Mo steel tested in the normalized and

tempered condition. Data are for test temperatures in the

range 482 to 704°C and for eight commercial heats.

P27-19
TEMPERATURE ( *C ) Figure 12. Estimated design allowable stresses as a function of temperature for modified 9 Cr-1 Mo steel. Design allowable

stress values for standard 9 Cr-1 Mo and 2 1/4 Cr-1 Mo steel

are also included for comparison.

Results of this study have shown that the thermal expansion

coefficient of modified 9 Cr-1 Mo steel is the same as that observed in

this study and reported in the literature for the standard 9 Cr-1 Mo

steel [4]. However, the thermal conductivity was found to be related to

the silicon content. Higher silicon content tended to be associated

with lower thermal conductivity. We have tested modified 9 Cr-1 Mo

heats containing 0.2 and 0.5% Si and recommend that the average values

for these two heats be used as typical values.

P27-20
I

6.2 Fatigue and Creep-Fatigue Tests

These tests are in progress at Oak Ridge National Laboratory,

Sandia National Laboratory, Northwestern University, and the University

of Connecticut. Data available have shown that the modified alloy has

about the same fatigue strength (total strain range versus number of

cycles to failure) as type 316 stainless steel in the range from 525 to

593°C. The alloy also has superior fatigue life beyond 105 cycles to

21/4 Cr-1 Mo, standard 9 Cr-1 Mo, and type 304 stainless steel.

Microstructural observations on fatigue-tested specimens are reported

by Jones [ 5]

6.3 Fatigue Crack Growth and Fracture Toughness

These tests are currently in progress at the Hanford Engineering

Development Laboratory. Results available showed that the crack growth

rate of the modified 9 Cr-1 Mo steel is about the same as that observed

for 21/4 Cr-1 Mo and 12 Cr-1 Mo steel in the range from room temperature to 538°C [6].

6.4 Steam and Air Oxidation

The behavior of modified 9 Cr-1 Mo steel and other alloys

(2 1/4 Cr-1 Mo, 9 Cr-2 Mo, 12 Cr-1 Mo, and type 304 stainless steel) in

superheated steam at 482 and 538°C are available for a period of

28,339 h. Results of this study have shown that silicon is very potent

in reducing the oxidation rate in steam at 538°C of chromium-molybdenum

steels [7]. Because of its lower silicon content, the modified alloy

showed slightly higher weight gain in steam at 538°C than that observed

for standard 9 Cr-1 Mo alloy.

Air oxidation data at 59 3° C were obtained on both the modified

9 Cr-1 Mo and 2 1/4 Cr-1 Mo steel for a period of 20,000 h. These data

showed that the weight gain for 21/4 Cr-1 Mo steel was about eight

times that observed for the modified 9 Cr-1 Mo alloy. Tests are now continuing to measure the weight gain for longer periods.

P27-21
6.5 Welding

The weldability of modified 9 Cr-1 Mo steel is being investigated

with the Varestraint hot cracking test, Battelle under-bead cracking

test, Tekken Y-groove test, hydrogen sensitivity test, and stress-relief

cracking test. Results available thus far have shown that the material

is free from hot cracking susceptibility, a preheat temperature of 200°C can prevent hydrogen sensitivity, and there are no stress-relief

cracking problems for a postweld heat-treatment temperature of 732°C.

The filler wire composition for gas tungsten arc welding, electrod com- position for shielded metal arc welding, and the combination of filler

wire and flux for submerged arc welding are currently being optimized to

satisfy the strength and ductility criteria set for the alloy.

7. Operating Experience on Modified 9 Cr-1 Mo Steel Tubes

Industrial operating experience, which aids earlier approval of the

data package for the ASME Code, is being obtained by installing tubes of

this alloy in various conventional power plants. The status of installation of modified 9 Cr-1 Mo tubes in various utility power plants is

summarized in Table 5. The range of utilities involved is international

United States, United Kingdom, and Canada. In most cases the tubes

being replaced are stainless steel. The longest operating time has been

reached for tubes installed at the Kingston Steam Plant. The tubes

installed in Kingston plant are shown in Fig. 13. The modified

9 Cr-1 Mo steel tubes were supplied with type 347 stainless steel safe

ends. The safe ends were welded on with Inconel 82 (ERNiCr3) filler

wire. The final welds between type 347 and type 321 (existing tubes) were made on location.

Tubes in the Agecroft power station went into operation in

April 1982 and those in Lambton and Nanticoke are expected to go into

operation in October 1982. This operating experience will be very useful for obtaining the approval of this alloy in the ASME Code.

8. Status of Commercialization

The use of a modified or new alloy in actual application requires

that it be included in ASTM specifications and various sections of the

P27-22
Table

5.

Current

status

of

testing

of

modified 9

Cr-1

Mo

steel

tubes

in

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P27-23 i
Modified 9 Cr-1 Mo tubes

Figure 13. The modified 9 Cr-1 Mo steel tubes before going into

operation in the TVA Kingston Steam power plant.

P27-24
ASME Boiler and Pressure Vessel Code . The ASTM specifications are in

fact prerequisites for the consideration of this alloy by the ASME Code

Committees. The schedule for the ASTM and ASME Code approval is as follows:

An application for an ASTM specification for the plate and tube

product was submitted for approval in May 1981. The specification

application for the forgings, piping, and fittings was submitted in

May 1982. The tube subcommittee has approved the specifications, which

now must be approved by the main committee. The plate specifications

are near approval, but the other specifications are still under con- sideration by the appropriate subcommittees.

The data package for the inclusion of modified 9 Cr-1 Mo steel in

Sect. I and VIII of the ASME Boiler and Pressure Vessel Code was submitted in June 1982. At the request of various users and producers, the

appropriate subcommittees of ASME are expected to start reviewing the

data package during September 1982.

8 . Summary

The status of development and commercialization of a modified

9 Cr-1 Mo alloy is presented. The alloy is modified by the additions

of niobium (0.06—0.10%) and vanadium ( 0. 1&—0. 25%) . The alloy is

recommended for use in the noramlized and tempered condition (1038°C for

1 h, air cooled to room temperature; 760°C for 1 h, air cooled to room

temperature). Heat-treatment, Charpy impact, tensile, and creep

properties of the alloy are discussed in detail and other properties are described briefly. Some of the key facts about modified 9 Cr-1 Mo alloy

include the following.

1. The ductile-to-brittle transition temperature (68-J) of

modified 9 Cr-1 Mo is lower and the upper-shelf energy higher than those

of standard 9 Cr-1 Mo alloy. These properties were much better in the

electroslag-remelted material.

2. The creep rupture strength of the modified alloy is higher than

that of the standard 9 Cr-1 Mo and 21/4 Cr-1 Mo steels for the entire

creep temperature range. The improvement in lC^-h creep-rupture strength

P27-25
is very significant at temperatures above 500°C. Higher strength is

also accompanied by total elongation and reduction of area values higher

than 15 and 70%, respectively.

3. Higher creep strength of the modified 9 Cr-1 Mo steel provides

design allowable stresses that exceed those of 2 1/4 Cr-1 Mo and

standard 9 Cr-1 Mo steel for the entire temperature range. At temperatures of at least 500°C, the design allowable stresses for the modified

alloy are twice those observed for 21/4 Cr-1 Mo and standard 9 Cr-1 Mo

steel.

4. The specifications for various products of this alloy have been

submitted to ASTM for approval. A data package has also been submitted

to the ASME Boiler and Pressure Vessel Code for approval of this

material in Sects. I and VIII.

5. Operating experience on this alloy is being obtained by

installing its tubes in various steam power plants.

The author thanks the following for their contributions:

R. E. McDonald (ORNL) and G. C. Bodine (CE, Chattanooga) for contributions to melting, fabrication, and heat treatment and helping to prepare

tubes for installation in various power plants; W. J. Stelzman (ORNL)

for Charpy V-notch studies; R. H. Baldwin (ORNL) for tensile and creep

testing; M. C. Cowgill (Westinghouse-Advanced Reactors Division) for

creep testing; J. F. King (ORNL) and C. D. Lundin (University of

Tennessee) for welding studies; R. K. Williams (ORNL) for physical

properties; and M. K. Booker (ORNL) for the data analysis. Our special

thanks are also due to the research staff of Climax Molybdenum Company

of Michigan for its contributions to the success of this program.

Continued encouragement and support of E. E. Hoffman (DOE, Oak Ridge

Operations) and P. Patriarca (ORNL) to the success of this program also

are greatly appreciated.

P27-26
9. References

[1] Sikka, V. K. , Ward, C. T., and Thomas, K. C. "Modified 9 Cr-1 Mo

Steel — An Improved Alloy for Steam Generator Application,"

presented at the ASM International Conference on Production,

Fabrication, Properties and Application of Ferritic Steels for

High-Temperature Applications, Warren, Pa., Oct. 6—8, 1981, to be

published.

[2] Bodine, G. C., and McDonald, R. E. "Laboratory and Pilot Commercial

Process/Product Development of Modified 9 Cr-1 Mo Ferritic Alloy,"

presented at the ASM International Conference on Production,

Fabrication, Properties, and Application of Ferritic Steels for

High-Temperature Applications, Warren, Pa., Oct. 6—8, 1981, to be

published.

[3] Orr, J. , et al. "The Physical Metallurgy of Chromium-Molybdenum

Steels for Fast Reactor Boilers," pp. 91—109 in M. Conf. Ferritic

Steels for Fast Reactor Steam Generators, 30 May—2 June 1977 , British Nuclear Energy Society, London.

[4] Williams, R. K. , et al. "The Physical Properties of 9 Cr-1 Mo

Steel from 300 to 1000 K," paper presented at the 17th Int. Thermal

Conductivity Conference, Gaithersburg, Md., June 15—19, 1981, to be

published in the proceedings.

[5] Jones, W. B. "Effects of Mechanical Cycling on the Substructure of

Modified 9 Cr-1 Mo Ferritic Steel," presented at the ASM

International Conference on Production, Fabrication, Properties and

Application of Ferritic Steels for High Temperature Applications,

Warren, Pa., Oct. 6—8, 1981, to be published.

[6] James, L. A. Fatigue-Crack Growth Behavior in Ferritic Alloys for

Potential GCFR Structural Applications , HEDL-TME 80-71, Hanford

Engineering Development Laboratory, Richland, Wash., December 1980.

[7] Griess, J. C., and Maxwell, W. A. The Long-Term Oxidation of

Selected Alloys in Superheated Steam at 482 and 538° C, ORNL-5771,

Oak Ridge National Laboratory, Oak Ridge, Tenn. , July 1981.

P27-27

S i Mo Ductile Iron for Elevated Temperature

Service to Conserve Chromium

Jay Janowak*

INTRODUCTION

High silicon molybdenum ductile cast iron is a relatively low cost

ferritic heat resistant material. In the trade it is called high

silicon moly iron or just plain SiMo. It is already in direct

competition with high chromium steels in some applications on a purely

economic basis. It is a popular material for turbo charger housings and

exhaust manifolds. It has been found to be the superior material for

rabble teeth in a moly-sulfide ore roaster. And finally, SiMo is a

candidate for furnace grates. All of these either are, or could become,

high chrome steel applications. Should there be a "critical" need to

reduce our use of chromium or to decrease our dependency on chromium,

we believe the high silicon moly ductile cast irons could be considered

as an option to high chromium steel. It is probable, but somewhat

speculative, that relatively minor design changes could allow SiMo

ductile iron to function in place of high chromium steels in numerous

situations. In the next few minutes four applications will be reviewed

to provide a feeling for the characteristic properties of this material.

SiMo ductile is basically a standard ferritic nodular (ASTM 60-40-18)

to which 1.5% silicon and 1% molybdenum are added to create the 4 silicon -

1 moly heat resistant ductile cast iron.

* Jay Janowak is Manager of Foundry Development, Climax Molybdenum Company

P. 0. Box 397> Arlington Heights, Illinois 60006. This is the oral

text of a workshop presentation "Trends in Critical Materials Requirements for Steels of the Future Conservation and Substitution Technol-

°qy for Chromium" at Vanderbilt University, October 5, 1982

P28-1
- 2

-

TURBOCHARGERS

The Si-Mo nodular iron concept was developed at the Climax

research laboratory in the sixties. Among early anticipated

applications were relatively low cost tubing for use in hydrogen gas

generators and forming dies for titanium. These did not materialize,

but Climax engineers brought the concept to the attention of a prominent

turbocharger manufacturer, and, after an extensive development program,

Si - 0.6% Mo ductile iron emerged as an improved housing material.

Figure 1 shows an automotive turbocharger. To give an idea of size, a

small car unit would fit inside a 1 gallon milk container. A cut away

view is shown in Figure 2.

Turbochargers are like miniature air compressors. Hot exhaust

gases spin the turbine which through a common shaft drives the blower.

On demand, fresh compressed air can be forced into the combustion

chamber of the engine to provide added power in trucks and off-road

vehicles. Using a different approach, this concept is applied to improve

auto fuel economy through utilization of smaller, lighter weight engines.

Exhaust gases are typically in the 1300 to 15oo°F range. Housings

can become red hot as shown in Figure 3- Dimensional stability , oxidation

resistance, and high temperature strength and toughness are important

properties provided by SiMo ductile iron in this application. Let me

emphasize the critical criteria.

Conta i nment - Housings must be able to contain a "wheel burst" for

obvious safety considerations. On vehicles, heavy walls are prohibitive

because of weight and space restrictions. Therefore, the relatively

P28-2
-3-

weaker gray irons are not practical. Consequently, the alternatives

are narrowed to ferritic or austenitic ductile irons or stainless steel.

However component complexity favors the castability of a ductile iron

over stainless steel options. While austenitic irons are stronger at

temperature, the ferritic irons have better ductility. Since containment

is related to a combination of strength and ductility, either iron will

serve the purpose.

Figure 4 shows strength and toughness at 1300°F of a 4% silicon

ductile iron with increasing molybdenum contents. Notice that molybdenum

substantially increases strength at 1300°F and also note the elongation

of over 50% at a level of \% molybdenum. Ductile behavior is evidenced

in Figure 5 which shows a housing that contained a wheel burst at temperature during a scheduled "run-away" test on a dynamometer.

Dimensional Stability - Dimensional stability is essential in general,

but, specifically, the turbine blade/housing "air-gap" must be maintained

at a minimum level. Divider wall stability is also very important.

Thermal expansion, phase changes, oxidation growth and distortion can

influence the "air-gap" and other contours visible in the cutaway view

(Figure 2). Silicon and molybdenum both work in the direction of improving the dimensional stability of nodular irons. Ferritic ductile iron

has a very low coefficient of thermal expansion which is dimensionally

beneficial compared to an austenitic material. The addition of silicon

further improves dimensional stability by raising the temperature of the

phase change to austenite as shown in Figure 6. Note that, at k% silicon,

ferrite is stable up to over 1500°F.

P28-3
-kMolybdenum provides dimensional stability by increasing the

strength at temperature and by improving creep resistance. In addition,

molybdenum adds resistance to stress relaxation, thereby reducing

warping and distortion. Figure 7 shows constrained thermal fatigue

specimens. Both are high silicon ferritic iron. One sample accumulated

a bulge - it had no molybdenum. The other sample contained .6% molybdenum. It showeo less bulge and survived 3 times the number of cycles

to failure. This benefit of molybdenum applies to both ferritic and

austenitic irons.

Oxidation - Ductile iron is inherently more oxidation resistant than

gray iron because the graphite is in discreet spheres as will be seen

later. This is a mechanical influence related to the transmission of

oxygen. Chemically, silicon and aluminum can add greatly to oxidation

resistance in cast irons. The silicon effect is shown in Figure 8.

Here, weight gain is plotted versus time at 1200°F for various silicon

levels. Notice the initial weight gain in all cases. This is explained

by the mechanism of oxidation resistance apparent in Figure 9 • Silicon

preferentially oxidizes at the surface to provide a silicon oxide layer

which prevents further penetration of oxygen. Notice also the discreet

spheres of graphite characteristic of ductile iron. This is a different

graphite shape than that of gray cast iron which has interconnected

graphite flakes that create a virtual "pipe line" for oxygen. This

silicon scale is particularly tenacious. Without adequate silicon, this

scale would exfoliate and rapidly abrade the turbine blades. Silicon

provides this benefit to both ferritic and austenitic irons. By limiting oxidation and thereby growth, high silicon irons have additional

dimensional stability.

P28-4
-5-

To bring this section to an end let me point out that turbocharger

housing materials can be classified according to service temperature.

Gray irons can serve up to 1000°F. Unalloyed ferritic nodular irons

can serve up to 1300°F, and with additions of silicon and molybdenum

up to 1500°F. Austenitic irons with silicon and molybdenum may serve

up to 1700°F while temperatures beyond that point may require high

chromium steels or superalloys. High chromium stainless steels are

alternatives to the irons at all these nominal temperature levels,

particularly when special thermal or environmental situations challenge

the cast irons. In such cases, silicon and molybdenum additions to

ductile iron make it a suitable alternative to high chromium steels.

EXHAUST MANIFOLDS

Exhaust manifolds are another hot engine item (Figure 10).

Traditionally, automotive exhaust manifolds were gray iron. However,

higher operating temperatures, demand for weight reduction, and the

desirable aspect of low thermal conductivity (for faster catalytic

converter light off) has resulted in a change to predominantly unalloyed

ferritic nodular iron. Most recently, a substantial interest has been

shown in fabricated stainless steel exhaust manifolds and there are

indications that some of the new cars may have these as standard. The

principle advantage of fabricated stainless steel is weight reduction -

since the remaining properties should be adequately provided by the SiMo

ductile irons. In fact, it is believed that SiMo ferritic ductile irons

are superior to the fabricated stainless steel in design flexibility,

noise damping and distortion resistance. Thermal distortion can make

it impossible to replace an exhaust manifold once removed from the car.

P28-5
- 6

-

In extreme cases, during operation, thermal distortion can lead to

exhaust gas leaks which can be a noise problem and a potential fire

hazard. Building a further case for the iron, it is important to

recognize that it is now possible, with new casting technology to

cast ductile iron exhaust manifolds having sections as thin as

one-sixteenth of an inch. The jury is still out on this application

but SiMo ductile irons are clearly in competition with stainless steel

having more than 12% chromium.

In a specific case involving very large diesel engines with

pistons over 15" in diameter, SiMo ductile iron replaced tubular

wrought 2iCr-1Mo steel exhaust manifolds. A design change was desired

by the manufacturer to reduce cost and improve performance. The new

design provided a common tuned exhaust system instead of multiple

individual exhaust tubes. Initially, unalloyed ferritic nodular iron

was considered and had the necessary or improved corrosion resistance

to condensate. However, concern over oxidation and creep led to the

selection of SiMo ductile iron which provided the necessary confidence

level for the new design. In this case, SiMo ductile iron replaced

2iCr-1Mo steel for a cost and performance benefit.

GRATES

Furnace grates (Figure 11) are another application where SiMo

ductile iron is being considered. An industrial incinerator operator

was looking for a lower cost alternative to HC steel. HC steel is

26-30%Cr, 4%Ni (max), .5%C(max), 2%Si(max), and 1%Mn(max). An enterprising foundry looking for new business decided to pursue the SiMo

iron with our guidance. A review of the application revealed certain

P28-6
- 7

-

needs. Metal to metal wear resistance was needed to protect against a

sliding action inherent in the design. Resistance to spike heating up

to 1500°F was desired, but normal grate operating temperatures were

said to be on the order of A00°F because the grates were effectively

insulated by the burning rubbish. In addition, the grates were cooled

by cold air forced up through the grates from below to provide oxygen

for burning. Toughness was important because motor blocks sometimes

were dropped from a height of 20 feet onto the grates. There was no

feeling for the corrosion resistance needed. A comparison of available

properties was made and shown in Figure 12.

Based on this information a decision was made "to go" and

trial castings were put into a controlled pilot experiment side by

side with HC steel. After 6 months the incinerator was shut down and

inspected. SiMo ductile grates showed less wear than HC steel. Then,

sometime after the 6 month check, the grates were found to have significantly deteriorated (Figures 13 and 1A). A preliminary investigation

revealed that molten aluminum had apparently fluxed the iron causing a

reduction in the melting point. The iron thus appeared melted in an

environment where temperatures only rarely reached 1500°F. The

investigation has begun only recently, and as of this moment, the

relative resistance of SiMo iron and HC steel to aluminum fluxing is

not known. In refuse systems where metal separation occurs after

burning, relative resistance to such fluxing action may end up being a

critical material selection criteria.

P28-7
- 8

-

Specific tests are planned to determine the relative resistance of

SiMo ductile and HC steel to aluminum fluxing. It is however clear from

the results to date that, except for this fluxing by aluminum, the SiMo

grates performed equally well or better than those of HC steel.

ROASTER TEETH - The last example is a case where SiMo ductile iron has

been found the superior material in an ore roasting furnace. Five years

ago, engineers at our moly sulfide roasting plants began a program to

improve the life of the roaster rabble teeth.

A number of rabble teeth standing upside down and waiting to be

put into service are shown in Figure 15* They hang from a beam that

rotates inside a roaster as shown in Figure 16. This view shows the

roaster in the empty and cold condition. The main shaft is also visible.

The tooth closest to the shaft is in location number one. In Figure 17

is a view showing the red hot rabble teeth doing their job which is to

roll over the molybdenum sulfide in an oxidizing atmosphere at up

to 1500°F. The rabble teeth are basically plows which must stand

up to an abrasive and especially hostile environment of hot molybdenum

oxide, oxygen, and high sulfur gases.

At times, during operation, it is necessary to maintain a free

flow condition. In Figure 18 is an operator poking to free teeth from

the substance being plowed when it becomes gummy creating excessive

drag on the shaft. You might say some impact resistance is needed.

The standard material was ferritic nodular iron with a silicon

carbide insert at the plow location for abrasion resistance. A number

P28-8
-9-

of materials were involved in the controlled experiment including

SiMo ductile iron, high chromium white irons and stainless steels.

The basis for selection of these materials for test was fairly

straightforward. The SiMo ductile iron was included on the premise

that it might last sufficiently longer than standard ferritic ductile

iron to meet the objective time and still be low cost. High chromium

white iron is well known for abrasion resistance but high temperature

endurance in this environment was unknown. Stainless steel was assumed

to be the best material to withstand the corrosive environment. It

was felt that significantly extended life could justify the added cost

of this high chromium-nickel option.

Results after one year of continuous service are apparent from

pictures of the teeth in Figures 19-22. Standard teeth in Figure 19 are

ferritic ductile iron with about 2 % silicon and no alloying elements.

The smallest tooth suffered the greatest deterioration and was closest

to the shaft where the environment is most hostile. The other teeth

were relatively further away from the shaft and in environments of

declining hostility. This same pattern follows with all of the materials

tested. CF8 (cast 304) stainless steel (l9%Cr~9%Ni) is shown in Figure 20,

28% chromium white iron in Figure 21, and 4 silicon - 1 molybdenum ductile

iron in Figure 22. Wear data are numerically summarized in Figure 23

in terms of "percent of original weight remaining" on the vertical axis

and roaster arm location along the horizontal axis. Position 1 is hotter

and more hostile. Position 14 is cooler and less hostile.

P28-9
- 10 -

The roaster teeth can be classified by performance into two

groups; the high silicon alloyed nodular iron (Alloys 1 and 8) - and all

the rest

; as clearly evident. There is a distinct difference in the way

the nodular irons eroded and the way the stainless steels and high Cr

white irons eroded. The nodular irons developed a thick oxide scale

which appeared to protect the metal surface and reduce the rate of

oxidation. The stainless steels, Alloys 3 and 9, did not develop a thick

protective scale, and thus, eroded very rapidly. Part of the rapid

deterioration of the stainless steel may have been related to the

nickel content. However a specific study would be necessary to verify

that possibility. The high Cr irons, Alloys 6 and 7» did develop an

oxide scale but it apparently had inadequate stability in the roaster

env i ronment

.

It is clear that the high chromium alloyed materials did not perform

any better than the standard alloy in the roaster environment. The

oxidation protection normally accompanying high chromium contents in

ferrous alloys is apparently just not effective in this environment.

On the other hand, it was certainly a pleasant surprise to find the

SiMo ductile iron as the most servicing material in this aggressive

env i ronment

.

In conclusion I would simply like to summarize these case histories

and restate our premise.

P28-10
- 11 -

1. Four silicon - one moly ductile cast iron is currently in

competition with high chromium steels for elevated temperature

service in exhaust system components such as tur bochdrger

housings and exhaust manifolds. SiMo is a common turbocharger

housing material. It has replaced 2£ Cr - 1 Mo in an exhaust

manifold. And, it is being challenged by fabricated stsir’ess

steel for automotive exhaust systems.

2. SiMo ductile iron is being considered as a replacement for

27 % chromium HC steel in industrial incinerator grates.

3. SiMo ductile iron has been found to be the superior rabble tooth

material in the highly corrosive - erosive environment of a moly

sulfide roaster over standard ferritic ductile iron, 2~J% chromium

white iron, and 19 % chromium - 3% nickel stainless steels.

A. It is probable, but speculative, with relatively minor changes i

design or performance requirements, that SiMo ductile cast iron

could replace high chromium steels in numerous high temperature

app 1 i cat i ons^ shou 1 d a critical need to do so develop.

P28-11
Figure 1 Automotive TurbochargerW

P28-12
3LM

TumOCH/WGtR

Figure 2 Automotive Turbocharger — Cutaway View

P28-13
Figure 3 Turbocharger in Service on Test StandL

P28-14
Figure 4 Tensile Properties of 4% Si Nodular

Cast Iron at 704 C (1300 F)

P28-15

REDUCTION

IN

AREA

AND

ELONGATION,

%
Figure 5 Turbocharger Housing Damage After Containment of Wheel Burst

P28-16
TRANSFORMATION

TEMPERATURE

,

(A,) C

TRANSFORMATION TEMPERATURE (A,)

Figure 6 Effect of Silicon Content on Acp of Nodular Irons at Various

Carbon Contents. Curves Are From the Literature, Data Points

Represent Climax Data (Dilatometry ) for 3.1%<C<3.8%.

P28-17

TEMPERATURE,
Figure 7 Constrained Thermal Fatigue Test

Nodular Irons. Sample on Left

Specimens of High-Silicon

Alloyed with 0.6% Mo.

P28-18
OXIDATION RESISTANCE

Figure 8 Effect of Silicon Content on Oxidation Resistance

of Nodular Iron at 650 C (1200 F)

P28-19
Figure 9 Photomicrograph of Oxidized High-Silicon Nodular Iron

Showing Adherent Silicon Oxide Scale

P28-20
Figure 10 Exhaust Manifold

P28-21
Figure 11 Commercial Incinerator Furnace Grate Before Service

P28-22
HC STEEL 4 Si-1 Mo

Tensile Strength

at Room Temperature 55,000 min. 80,000 min.

% Elongation

at Room Temperature 2-19 8-19

Brinell Hardness No. 190-223 190-240

Load Bearing Service up to 1200°F up to 1500°F

Rupture Stress

at 1300°F — 3300 psi

at 1400°F 2300 psi —

Figure 12 Comparison of Properties of High-Chromium (HC) Steel

and 4Si-lMo Nodular Iron

P28-23
Figure 13 Si-Mo Nodular Iron Furnace Grate Showing Damage Resulting

(Apparently) from Exposure to Molten Aluminum

P28-24
Figure 14 Si-Mo Nodular Iron Furnace Grate Showing Damage Resulting

(Apparently) from Exposure to Molten Aluminum

P28-25
Figure 15 Roaster Teeth Prior to Service

P28-26
Figure 16 Roaster Teeth in Position on Rabble Arm

P28-27
Figure 17 Roaster Teeth in Service

P28-28
«

I

Figure 18 Operator Using Steel Rod to Clean Roaster Teeth

P28-29

**rd*9L

**
UNALLOYED (2% Si) NODULAR IRON RABBLE TEETH

Figure 19 Unalloyed Nodular Iron Roaster Teeth

after One Year of Service

P28-30
CF8 (CAST 304) STAINLESS STEEL RABBLE TEETH

Figure 20 Cast 304 Stainless Roaster Teeth

after One Year of Service

P28-31
28% Cr WHITE IRON RABBLE TEETH

Figure 21 High-Chromium White Cast Iron Roaster Teeth

after One Year of Service

P28-32
4% Si - 1% Mo NODULAR IRON RABBLE TEETH

Figure 22 4Si-lMo Nodular Iron Roaster Teeth

after One Year of Service

P28-33
o

100

90

80

70

60

50

40

30

20

10 “

0

Shh\FJ

FLOW OF

MATERIAL

-

_L

ALLOY 6 (2 7C-27Cr WHITE IRON)

a ALLOY 7 (1 25C-27Cr WHITE IRON)

• ALLOY 8 (4. 5Si - 1 Mo NODULAR IRON) - DALLOY 9 (0 2C-19Cr-9Ni STEEL)

OSTANDARD ALLOY (FERRITIC NODULAR

IRON)

I I I

i i i i i 4 5 6 7 8 9 10 11 ROASTER ARM POSITION

12 13 14.WALL

Figure 23 Performance of Various Materials as Roaster Teeth in Climax Test

P28-34
FASRAD #82-12

THE 1982 STATUS REPORT ON FE-MN-AL STEELS*

Samir K. Banerji

Foote Mineral Company

Exton, PA 19341

INTRODUCTION

The development of the Fe-Mn-Al alloy system as a possible substitute

for the traditional Ni-Cr austenitic stainless steels in some applications

appears to be gaining some momentum. In the June 1981 Workshop on the

"Conservation and Substitution Technology for Critical Materials" held at

Vanderbilt University, Nashville, Tennessee, the author presented an update

of activities on this subject.

^The available literature until that time

was reviewed in that report and a brief summary of some preliminary work

(2 3) spurred by the two earlier papers of the author * was presented. Some

of those efforts have been completed, some have been forgotten, some are

still continuing and some new efforts have been started. It is particularly

gratifying to see that the results of some of these efforts are beginning

to appear in the literature.

(4) A paper on the oxidation resistance of Fe-Mn-Al-C alloys appeared

recently and another one on Mn-Al-Si austenitic steels for propellers of

sea-going vessels

^will be forthcoming shortly in the Metal Progress .

A third paper on the marine corrosion behavior of several Fe-Mn-Al

austenitic steels

^will be submitted for publication in the near future.

Another paper outlining the various properties and characteristics of the

Al-Mn-C austenitic steels has been submitted for publication in the Journal

of Metals . ^

Six papers were presented on this subject at the XXXVI Bra-

(8—13) zilian Society for Metals (ABM) Annual Congress in July 1981. These

*To be presented at the Public Workshop on "Trends in Critical Materials

Requirements for Steels of the Future - Conservation and Substitution

Technology for Chromium," Vanderbilt University, Nashville, Tennessee,

October 4-7, 1982.

P29-1
2

papers provided some preliminary results of a large-scale study undertaken

at various universities in cooperation with the Technological Center

Foundation of Minas Gerais-CETEC in Brazil. Two papers appeared

recently at the 35th Annual Conference of the Australasian Institute of

Metals dealing with the oxidation and sulfidation resistance of Fe-Mn-Al

type alloys. Brief summaries of these and the other ongoing work are

presented in this report.

SUMMARY OF THE NEW LITERATURE SINCE THE JUNE 1981 WORKSHOP

1. development of Oxidation Resistant Fe-Mn-Al Alloys" by J. C. Garcia

et. al. (ref. 4) .

The microstructure and oxidation behavior at 500° and 700°C (930° and

1290°F) of several Fe-Mn-Al-Si-C-X alloys, shown in Table 1, were studied.

Alloys A and B showed a dual-phase structure, predominantly austenitic

with different amounts of an ordered b.c.c. phase, when homogenized at

1100° and 1000°C, respectively. Alloys C and D were single phase austenitic

when quenched from 1100° and 1000 °C, respectively. Alloys E and F with Cu

additions were primarily ordered b.c.c. with the precipitation of a disordered f.c.c. phase. The oxidation behavior of these alloys were

measured by weight gain experiments in flowing oxygen (1 atm.) at 700° and

500°C for times up to 24 hours.

Table 1. Composition of six alloys studied

by Garcia et. al. (ref. 4).

Alloy Composition, wt %

Al Mn Si C Cu P S Fe

A 10.84 31.92 3.01 0.74 — 0.05 0.02 bal

B 9.05 33.30 1.41 0.88 — 0.02 0.02 bal

C 8.63 28.60 2.68 0.69 — 0.05 0.02 bal

D 7.25 24.04 2.06 1.01 — 0.02 0.02 bal

E 12.40 4.13 2.12 0.03 0.91 0.02 0.01 bal

F 12.05 4.20 1.93 0.03 8.90 0.02 0.01 bal

P29-2
3

Figure 1 shows that alloy E behaves similarly to type 304 stainless

steel at 700°C while the behavior of the rest of the alloys in 24 h runs

is far superior for the Fe-Al-Mn alloys than for a plain carbon steel.

Results reported in Fig. 2 indicate that Fe-Al-Mn alloys behave comparatively and superior to type 304 in 24 h experiments at 500°C. These

results are in agreement with other reports. The formation of

aluminum-rich surface oxide scales was a common observation in these

alloys during the oxidation tests. Addition of Cu appeared beneficial

for oxidation resistance. Finally, alloys A, B, C, and D were hot and

cold rolled at room temperature into 1 mm (0.04 in.) thick sheet without

intermediate anneals. Alloys E and F became brittle at room temperature

probably because of brittleness associated with the ordered parent phase.

2 . "An Austenitic Mn-Al-Si Steel for Propellers of Sea-Going Vessels " by R. Wang and F. H. Beck (ref. 5) .

This soon to be published paper describes the selection of an Fe-MnA1 steel for the propeller of a sea-going fishing vessel in People's

Republic of China, launched on August 10, 1980, and still sailing safely.

The mechanical properties and marine corrosion resistance of as-cast

alloys of Table 2 were studied and compared with Mn-Fe brass which is the

traditional propeller material and austenitic stainless steel. The

strength and hardness of the Fe-Mn-Al alloys were higher than those of the

brass. These alloys were approximately 20% lighter than the brass and

also lighter than the stainless steel.

The marine corrosion tests were first carried out by Immersing

samples in boiling seawater for 118 hours. Alloy no. 1 and the Ni-Cr

stainless steel showed no evidence of rusting. Alloy no. 3 had a small

rust spot but alloys 2 and 4 were obviously rusty. Similar immersion

P29-3
4

tests lasting 768 hours (32 days) were also carried out at room temperature.

Alloys 2 and 4 were covered by a brown rust film, alloy 3 showed a thin

yellowish film, the Ni-Cr stainless had a little track of rust, but alloy

no. 1 was free from any obvious rust. Based on this the alloy no. 1 was

chosen for the propeller of the fishing vessel mentioned earlier.

Subsequently, the polarization curves in artificial seawater were

determined for the alloy no. 1 and the Ni-Cr stainless steel. Interestingly

enough, the corrosion current density (I corr) and the corrosion rate (Rcorr ) 3

were nearly identical, being 2.905x10 NA/Cm and 1.2 mpy, respectively.

Table 2. Compositions of Alloys Tested by Wang and Beck (ref. 5).

No. Code Name %C %Si %Mn %A1 %P %S %RE

1 ^O^IO81 1.01 1.31 30.5 10.4 0.007 0.006 0.056

2 M1120AI

7 0.50 0.30 22.3 7.2 0.005 0.003 0.038

3 Mn^QAly Si 0.23 1.25 29.2 7.4 0.008 0.003 0.061

4 ^20^5 0.13 0.28 20.7 5.6 0.007 0.005 0.063

The Main Composition of the Compared Materials

%C %Cr %Ni %Ti %Cu %Zn %Ma %Fe

5 Cr^gNiqTi 0.08 18.59 9.63 1.07

6 Mn-Fe Brass 55 41 3 1

55-3-1

3 . "Marine Corrosion Behavior of Several Fe-Mn-Al Austenitic Steels" by Wang and Rapp (ref. 6 ) .

This very recent work at the Ohio State University has concentrated

on the electrochemical corrosion behavior of several Fe-Mn-Al alloys.

The potentiodynamic polarization tests in artificial seawater or in

INH2 SO4 were performed on seven different alloys of varying compositions

and the results were compared with two standard austenitic stainless steels

P29-4
5

of types 316 and 321. The composition of the various alloys tested and the

results of the corrosion tests are shown in Table 3. Alloy no. 3

(30Mn-10Al-lC) showed the best corrosion resistance of all the seven alloys

tested with a linear corrosion rate somewhere in between types 316 and 321

stainless. Additionally, this alloy in INH2 SO4 also showed a passive region

within the potential of +0.5V to +1.5V.

The shape of the polarization curves for all these alloys in artificial

seawater including the 316 and 321 stainless steels were nearly the same,

all lacking an obvious passive region. However, with increasing A1 and/or

Si and decreasing C contents, there is a more stable region at anodic

potential near Ecorr . The shapes of the polarization curves for alloy

no. 3 and 316 stainless were nearly identical. Based on this and the results

in Table 3, it was concluded that A1 and Si are beneficial in the marine

corrosion resistance of these alloys. The carbon content should be kept

to a minimum necessary to maintain an austenitic structure.*

4. "An Evaluation of Al-Mn-C Austenitic Steel" by Rao et. al. (ref. 7) .

The mechanical properties and preliminary oxidation and corrosion

behavior of a 32Mn-9.6Al-1.0C-Fe alloy were studied. The alloy was found

to be age-hardenable with the peak hardening occuring around 600 °C aging

temperature. The mechanical properties of the steel in this peak-aged

condition were: yield strength = 90 kg/mm (128 ksi) , ultimate tensile

strength = 100 kg/mm^ (142 ksi) , and 40% elongation. The impact properties

in the solution annealed condition were good, but deteriorated drastically

on peak aging. The alloy exhibited excellent formability characteristics,

comparable to 304 stainless steel in 1.45 and 0.63 mm gage thicknesses.

The oxidation resistance in 100 hour tests in air showed the alloy to be

only marginally inferior to 304 SS at 700°C. This is in agreement with the

*Further work is continuing at the Ohio State University by these authors.

P29-5
Table

3.

Chemical

Composition

and

Corrosion

Data

of

the

Samples

Tested

by

Wang

and

Rapp-

(ref.

6).

M

ua

m 9 . a) ao CM Ov 00 n* cn sf o m

0)

u 9 tn o cn I—(

00 oo rx

m rH c M U — o o © rH 00 o O o fH

•H

O 03

g • • • • • • • • rJ O OS B o o o o o o o o o U /“X 9 • >v Ov cn in in o rH xd- CM Mr 9 u <u a cn •n cn vO xO *d- cn o rH C U 4J e cn CM fx. ,«d- 00 CM Ov o in

xH

O (JV • • • • • • • • • rJ U OS o fH o rx -O cn CM CM •d- cn vi-KNI •h a cn sr fx* rx vO oo CM CM O n u 00 00 rH CM CM CM CM CM xd- c — • • • • • • • . •

01 00 a

fX rx px Px px px Px Px Px • 4J • > J1 3 00 CM fx. xO xd- -d- rx rx CM rl 60 O' 00 o CM 00 00 00 cn cn xd- 3 -H W • # • • • . • • • V 9 ^x m vO cn cn cn cn cn cn in W » (30 CM CM CM CM CM CM CM CM CM

4JC CM • at wa

m cn o o o U h U U ao O xfl- CM *d- cn 00 vO ?H u U 0"x rx CT\ rx m px rx ov px rH o 3 9< U H 3 • • • • • • . • • • o o CM rH px rH px vO sr O rH 00 ^H rH

CO •H xd" CM CM px ov ov vO

O 4J rH X» iH CM vO rx m CM • C h xd- x3- m VO vO vO vO V0 n*. u 0) u • • • • • • • • • u 4J Oz-N O O o o o O © O o o O CJ> 1 1 i 1 1 1 1 o Oh M

CM Z • O

fx. •H o H • rH O 3 xd U • om o i—i

as • CM

K xO- cn

Wt N1

00• o

vO• Ov * iH co

•H rx Ov U cn m • • m ° 00 o ^H rH aio NO CM cn CM CM

m xd- 00

U iH rH CM rH o »d- in o

CO • « • • • • • • 3 o o o o o rH rH o

•H

CO X

m Ov cn 00 Px CM 00

rH rx cn cn CM cn cn VO < • • • • • . • ov oo 00 00 00 00 vO

CM in o o rH o CM e

VO rH vO rH 00

m rH cn • . • • • • « • X rH o o o ov o o © cn cn cn CM cn cn cn mCM 00 rH m xd- CM m rx vO u o o o o CM m o O o • • • . • • • o o rH rH rH o rH rH o

01 0)

ucn1 rH

rH1 <•1 CM •3 B 1 rH ® ' (ft o 1 00 O | rH O CO rH >d- CM rH xd- 1 o z VO iH 1 ^

rH

O 1 CM 1 | CM CM

«—1

CM rH 1 o rH 1 1 cn cn rH W rH ao rH rH 00 CM o *—i

CM cn >dm vO oo OV z 6

P29-6
7

previous work. At 900°C the oxidation resistance was considerably

inferior, again confirming the results of other workers. The corrosion

rate of half immersed samples in 7% NaCl at 60°C was nearly twice that of

304 SS. In 50% HC1 at room temperature the results were considerably poorer.

This work was summarized in my earlier report.

Additional work on the phase stability and transformation characteristics, microstructural correlation, and kinetics of aging are being completed.

In view of the higher stability of the austenite on deformation or cooling

to liquid nitrogen, these alloys do seem to have a good potential in cryogenic applications.

5. "The Resistance to Oxidation and the Electrical Resistivity of an Fe-Mn-Al System Alloy" by Casteletti and Spinelli (ref. 8) .

An alloy with 31.22% Mn, 7.54% Al, 1.34% Si, and 0.93% C (pouring 13)

and a non-silicon version of this alloy were tested. Oxidation tests were

carried out in air at 700° and 800°C for times up to 70 hours and compared

with Type 304 stainless steel. The resistivity of the alloy (pouring 13)

was measured over the temperature range of 80°-812°C using 80 mm long wires

with 2 mm4 cross-sectional area. Figure 3 shows that the oxidation resistance of the alloy at 700°C is comparable to 304 SS. This is in agreement

with previous work^*^’^ At 800°C the oxidation resistance is inferior to

that of 304 SS, Fig. 4. Nevertheless, the oxide layer did stabilize after

about 40 hours at 800°C. The alloy without Si had a substantially inferior

oxidation resistance at both temperatures. The electrical resistivity

values shown in Fig. 5 are higher than those of commercial Ni-Cr alloys and

Kanthal

.

Based on these results it was concluded that these alloys represented

a promising option in replacing the heat resistant alloys of Ni-Cr, Fe-Ni-Cr,

and Fe-Cr-Al at temperatures up to 800° C.

P29-7
8

6 . " Study of Oxidation of an Fe-Al-Mn-C Alloy at High Temperatures" by

Assuncao et. al. (ref. 9) .

A low carbon alloy (36.5 Mn, 7.4 Al, 0.13 C, bal. Fe) was studied in

this work. The oxidation behavior at 1000° and 1100°C was investigated

in air, pure oxygen and nitrogen for times up to 8 hours (50 hours in one

case) . The results of the oxidation tests in air do not indicate good

performance at 1000°C and 1100°C. This is in agreement with previous

(7 8) results. * The outer oxide layer which spalled-off was identified to

be Mn0 .Fe 203 which is consistent with the work of Rao et. al.^ The

inner adhering oxide layer was rich in Fe203 followed by a region of an

aluminum-rich precipitate, the existence of which was correlated with the

presence of nitrogen in the oxidizing gas atmosphere, suggesting the

precipitate to be nitride or carbonitride complexes of Al and Fe. The

oxidation resistance of the alloy in pure oxygen was excellent at 1000°C.

7. " Study of the Microstructural Transformations in TIG Welding of an Fe-Al-Mn Alloy " by Bushinelli et. al. (ref. 10) .

Again, a low carbon alloy was used with a chemistry nearly identical

(9) to the previous study. 3 mm thick cold-rolled plates in homogenized

(1000°C/2h, water quenched) and cold-worked condition were used for the

welding test. Due to the low carbon content (.12%), the alloy in both

these conditions showed a duplex microstructure consisting of austenite

and 6-ferrite, but the fraction of the ferrite phase was smaller and

spheroidized in the homogenized condition. The TIG welding was selected

because this is virtually an inconsumable electrode process. The welding

parameters were: 105-110A, 14-16V, weld rate of 20 cm/min. , and a pass

length of 75 cm.

The transformation of 6-ferrite in the weld heat-affected-zone (HAZ)

of the cold-rolled plates caused embrittlement and the fracture to occur

P29-8
9

in the HAZ at stress values about half the strength of the base metal. The

tensile tests of the welded plates in the homogenized condition were invalid

due to lack of weld penetration. The composition of the 6-ferrite in these

alloys was richer in A1 and poorer in Mn. Homogenization reduced the A1

content of the ferrite. The transformed 6-ferrite regions in the HAZ of

the cold-worked plates were also rich in A1 and low in Mn compared to the

austenitic matrix. Therefore, it was believed that the embrittlement in

this case occurred due to the precipitation of an intermetallic (Fe 3Al

type) or carbidic phase in the transformed ferrite regions rather than

3-manganese formation.

The homogenized plates did not show any transformation of the ferrite

in the HAZ. The HAZ microstructure remained unchanged from the base metal

microstructure of the homogenized plates, i.e., spheroidized ferrite with

smaller volume fraction and reduced A1 content. Although the results of

the tensile test in this case were invalid, it is believed that the weld

embrittlement could be eliminated by the prior homogenization treatment.

However, the ideal solution would be to minimize or eliminate the ferrite

phase altogether by increasing the austenite stabilizing element. In this

case a higher carbon version of the alloy could be useful.

8 . "Mechanical Properties of an Austenitic Steel of the System Fe-Mn-Al" by Casteletti and Spinelli (ref. 11) .

The same alloy (pouring 13) used in the other study of these

(8) authors, 3IMn-7.5Al-1.3Si-0.93C-Fe, was used for this work as well.

Tensile properties were measured in the temperature range of 24°-815°C

and also after various levels of cold-reductions ranging from 0 to 80%.

Low-temperature impact strength was measured over a temperature range

of +20°C to -196°C. The microstructure of the as-cast alloy showed

P29-9
10

large austenitic grains with some dendritic structure in the grain interiors.

After forging (at 900°C) and solution annealing at 1050°C followed by oil

quenching, a typical annealed austenitic microstructure was obtained.

In Table 4 are presented the results of the tensile tests in the ascast and forged states of the alloy. There is an increase of approximately

35% in the tensile strength and 40% in the yield strength of the worked

alloy with practically no loss of ductility compared to the as-cast alloy.

The mechanical properties of the as-cast alloy are still better than most

standard austenitic stainless steels.

Table 4. Results of the Tensile Strength Tests

for the As-Cast and Forged Material (ref. 11)

Material

Tensile

Strength

MN .m

- ^

0.2% Yield

Strength

MN.m^

Elongation

%

Reduction

in Area

%

As-Cast 738 401 50 —

Forged 997 567 48 59

Figure 6 shows the temperature dependence of the tensile strength of

the forged alloy in the temperature range of 24-815 °C. Note that over the

temperature range 24°-600°C, the tensile strength values for the alloy

under study are 30% higher than those of 202 SS; 50% higher than those

of 316 SS; and 68% higher than those of 304 SS.

Figure 7 shows the yield strength as a function of temperature for

the alloy studied compared to the standard stainless steels. It can be

seen that the yield strength values for the Mn-Al steel are on average

87% higher than those for the other alloys, in the entire range of

temperatures studied.

P29-10
11

One can also see a slight decline in the yield strength values in the

temperature range of 150-300°C, where there occurs the phenomenon of dynamic

aging of the alloy. Between the temperatures of 400° and 550°C, the yield

strength values exhibit a peak, possibly indicating the occurrence of

hardening phenomena due to precipitation in this range of temperatures.

Figures 8 and 9 show the dependence of tensile and yield strengths

on the percentage of cold reduction. Note that the rate of increase in

tensile strength with percent reduction is about the same for the Mn-Al

steel (pouring 13) as the standard stainless steels. However, the absolute

values are considerably higher due to the initially high tensile strength

of the alloy. On the other hand, the yield strength initially increases

more rapidly than the other alloys (Figure 9). As expected, the ductility

decreases as strength increases with cold reduction, see Figure 10. The

rate of drop in ductility is approximately the same as the standard stainless steels, although the absolute values are lower due to higher strength

of the Mn-Al alloy.

Figure 11 shows that the strain-hardening response of this alloy is

lower than those of two other stainless steels — Types 201 and 301.

Figure 12 compares the impact strength of the Fe-Mn-Al alloy with several

other steels from room temperature down to liquid nitrogen temperature. The

ductile-brittle transition temperature of this alloy in the as-forged condition is about -80°C and comparable to the 310 SS.

The specific gravity of the alloy was 6.85 g/cm^. This coupled with

the high strength levels provides a very attractive strength-to-weight

ratio in this alloy.

P29-11
12

9 . "Precipitation Reactions In Low-Carbon Fe-Mn-Al Alloys" by Tschlptschln

and Goldenstein (ref. 12) .

Two low-carbon alloys with .13C, 36.5Mn, 7A1 were studied; one without

Si (#T-04) and one with 1% Si (/AT—05) . After homogenizing at 1000°C for

times up to 24 hours, the microstructure still consisted of islands of

delta ferrite in an austenite matrix. The alloys were aged at 600°-800°C

for times ranging from 15 secs, to 200 hours. The transformation of

6-ferrite into lamellar carbide + ferrite was observed as a result of

aging. The transformed ferrite was an ordered phase of the Fe3Al type and

caused embrittlement of the alloy. The transformation nucleated at the

austenite/6-ferrite boundaries and the interlamellar spacing increased

with temperature.

No significant changes in hardness were observed as a result of

aging, as shown in Table 5. This lack of age-hardening is contrary to

other observations, probably because of the low carbon content of

these two alloys. At 1000°C both the alloys developed thick oxide layers

(9) when heated in air as observed by other workers.

Table 5. Vickers Hardness of Alloys T-04 and T-05 (ref. 12)

Heat Treatments: Carried[ Out

1000°C — 2 h

1000°C - 2 h & 800°C - 1 h

1000°C - 2 h & 800°C - 3 h

1000°C - 2 h & 700°C - 1 h

1000°C - 2 h & 700°C - 3 h

1000°C - 2 h & 600°C - 15 s 1000°C - 2 h & 600°C - 30 s 1000°C - 2 h & 600°C - 2 m

1000°C - 2 h 4 600°C - 10 m

1000 °c - 2 h & 600°C - 30 m

1000°C - 2 h 4 600°C - 1 h

1000°C - 2 h & 600 °C - 3 h

1000°C - 2 h & 635°C - 200 h

VH Hardness

Alloy T-04 Alloy T-05

179 163

158 158

165 157

160 165

161 171

152 159

170 155

148 154

157 164

148 162

160 145

143 153

179 175

P29-12
13

10 . MDomain of the Austenitic Phase in the Fe-Mn-Al System at 1000°C"

by Branco and Boratto (ref. 13) .

The 1000°C isothermal section of the carbon-free Fe-Mn-Al ternary

system was determined to establish the austenitic phase field. The

samples were held at 1000°C for 168 hours under 1 atm. argon pressure

before quenching in water. X-ray diffractometry , optical microscopy and

micro-hardness measurements were used to identify the various phases.

The results showed a slightly more limited austenitic phase field than

that proposed earlier. The maximum aluminum solubility in the

austenite at 1000°C was 15 atomic percent in the carbon-free ternary

system.

11. "Elements Essential to High Temperature Sulphidation Resistant IronBased Alloys" by Tomas (ref. 14) .

Preliminary results of this work on Femnal 30 alloy (30.4Mn, 7.6A1,

0.8C, bal. Fe) was summarized in my report last year.^ This and some

additional work is now completed. The purpose was to develop iron-based

alloys free of chromium to resist sulphidation attack at high temperatures

encountered in materials used for fossil fuel and coal conversion plants

due to the presence of sulfur.

A large number of alloys were tested for sulphidation resistance in a

reducing atmosphere of H2+ 1 vol.% H2 S for 9 hours at 700° and 800°C. The

weight gain data^^ shown in Figure 13 exhibited the best performance by

pure Mo and two iron-based alloys containing 5% A1 and 7% (Fenbal 7) or

17% (Fenbal 17) niobium. The Femnal 30 and a similar Si-bearing alloy

#11-10-4B (30.5Mn, 8.4A1, 1.5Si, 1.05C, bal. Fe) , supplied by me to

Dr. Tomas, also showed very good performance better than all the standard

Ni-Cr stainless and high Mo or Nb alloys.

P29-13
14

In both these alloys a relatively thin adherent sulfide scale formed,

see for example Figure 14 showing the scale formed in Femnal 30 after 9

hours at 700°C. The scale was duplex in nature, the outer layer being

predominantly MnS with very low iron content. Figure 15(a). Underneath

this outer layer of MnS there forms a thin layer of niddle-like aluminum

sulfides, either AI2

S 3 or FeAl2 S 4 , protruding into the inner matrix

(Figure 14) . This provides excellent scale/matrix adhesion and spallation

resistance to the scale. The slightly inferior result on the 11-10-4B

alloy (Figure 13) was believed to be due to the presence of iron-rich

sulfide phase in the outer layer of the scale, the lighter phase in

Figure 15(b). This could be due to the presence of Si in this alloy.

Apart from this, the scale on both this and the Femnal 30 alloys were

similar in nature.

British Gas Corporation has tested both Femnal 30 and the 11-10-4B

alloy in an oil gassifier test rig. The conditions were thought to be

reducing, containing about 0. 5-1.0 vol. % H2 S. However, both alloys

failed catastrophically after 300 hours at 750°C. Analysis of the

reaction products revealed oxides and a liquid slag. This indicated

that more oxygen was present than previously assumed. It has been shown

earlier that such alloys do not resist oxidation/sulphidation attack

(18) in multicomponent gas mixtures. However, in reducing/sulphidizing

atmospheres these alloys offer excellent potential at least up to

700°-800°C. At 700°C in this atmosphere these alloys are at least ten

times more resistant than the 300 series stainless steels.

P29-14
15

12 . "The Development of Fe-Al Alloys for High Temperature Oxidation

Resistance" by Jackson and Wallwork (ref. 15) .

This work deals with the carbon-free, low-Mn, ferritic alloys of the

Fe-Mn-Al system in an attempt to improve oxidation resistance at the higher

temperature range of 800°-1000°C. The high-manganese austenitic alloys do

not appear to have suitable oxidation resistance at these higher temperatures as discussed earlier. ^*^*9) Approximately fifty Fe-Mn-Al and several

Fe-Mn-Cr-Al alloys were prepared by a levitation melting technique. Specimens in the as-cast condition were oxidized at 800°C for 24 hours under an

oxygen pressure of 200 torr. Some selected alloys were also tested for

100 hours at 600°, 800°, and 1000°C.

Table 6 shows the weight gain data in 800°C/24 hour tests on the

various Fe-Mn-Al alloys. Based on these results the most optimum composition for 800°C service was: 5-10 Mn, 6-10 Al, bal. Fe.

Table 6. Weight Gains for 24 Hours Oxidation (AW/24)

at 800°C (mg. cm. ”2). Marked Area Showing

Optimum Composition Range for Fe-Mn-Al

System (ref. 15).

10

11

12

13

14

15

16

,062 .061

.053

.091

.155 .138

.166

.136

.074

,155

151 .191

*Specimen only oxidized for 18 hours.

P29-15
16

The weight gains and compositions of some of the Fe-Mn-Cr-Al alloys

oxidized for 24 and 100 hours at 800°C under 200 torr oxygen pressure are

shown in Table 7. The most protective scales with minimum weight gain

were found in alloys containing 5% Al, <15% Mn, and >5% Cr.

Table 7. Weight Gain and Composition of Some

Fe-Mn-Cr-Al Alloys (ref. 15).

Alloy

Composition Weight gain at 800 °C (mg cm Fe-%Mn-%Al-%Cr 24 Hours 100 hours

5 4 5 .292 (18 hrs) —

5 5 5 .068 .327

7.5 4 5 .196 .420

7.5 5 3 .122 .493

7.5 5 5 .063 .215

10 4 5 .115 —

10 4 12 .061 —

15 5 5 2.663 —

1 two alloys tested for 100 hours at 600° , 800°, and

• oxygen pressure are shown in Table 8 . Convoluted

with almost identical morphologies and composed of an outer layer rich in

Mn and Al, and an inner layer of AI2O 3 , formed on both alloys oxidized at

1000°C.

Table 8 . 100 Hour Oxidation Tests (ref. 15)

Alloy

Composition

Weight Gain

(mg cm"2

)

Fe-%Mn-%Al 600°C 800°C 1000°C

10 6 .171 .117 .450

20 8 .101 .242* .253*

*Weight gains unreliable due to oxide spallation.

P29-16
17

Based on this work the authors concluded that suitable alloys from

the Fe-Mn-Al system can be produced with oxidation resistance at 800°-1000°C

equivalent to that of heat resisting stainless steels. The final formulation of these alloys by possible addition of Cr, C, and rare earths is now

being undertaken in order to develop better room temperature properties and

high temperature strength.

OTHER UNPUBLISHED AND ONGOING WORK

1. University of Leeds, U.K. - G. E. Hale (ref. 19)

This work performed by G. E. Hale constitutes a part of his soon to

be submitted Ph.D. thesis at the University of Leeds. It concentrated

mainly on the austenitic alloy with about 30Mn, 8A1, 0.8-1.0C, bal. Fe.

Some work was also done on a duplex austenitic/ferritic alloy and a fully

ferritic material The composition of various alloys studied are shown

in Table 9.

Table 9. Alloy Compositions (ref. 19)

ALLOY

wn

ROOM

TTnyfOrD A 'TTTDTT COMPOSITION

nu • i£«rirJLr\Ai UxUL MICROSTRUCTURE Mn A1 C Fe Others

2 Austenitic 30.7.6 0.81 BAL 0.01 AN, 0.36Si

0.009S, 0.017P

3 Austenitic/

Ferritic

17.1 5.0 0.07 BAL 0.06Si

A Ferritic 8.9 6.9 0.07 BAL 0.007P

5 Austenitic 30.9 8.0 0.97 BAL 0.008N, 0.17Si

0.011S, 0.007P

P29-17
18

The Ferritic Alloy ;

The ferritic alloy No. 4 shows a distinct "script" phase together

with intergranular grain boundary precipitation in the as-cast state

(Figure 16) , but this can be removed by hot rolling and/or solution treatment at temperatures of 1100°C and above, followed by water quenching. The

exact nature of this "script" phase has not yet been determined. Unfortunately, this alloy has proved extremely difficult to machine and very

brittle at room temperature with a Charpy impact energy of only about

4 Joules. Aging after a high temperature solution treatment followed by

quenching results in a fine precipitation in both the grain boundaries and

grain interiors. No distinct increase in hardness was detected with

increasing aging time, however.

The Duplex Alloy :

The duplex alloy No. 3 shows some interesting effects and may well

be worth more detailed study. After hot rolling, the alloy contained 73%

austenite and 27% ferrite (the relative proportions of the two phases were

measured with an X-ray diffractometer) , but on solution treating at 1200°C

and higher followed by rapid quenching (preferably into iced brine) , the

alloy becomes single phase ferrite in fairly thin sections (1-2 mm)

Larger sections and/or slower cooling rates lead to austenite precipitation

both at ferrite grain boundaries and also within the grains.

Reheating the quenched material allows austenite reversion to occur

and the amount of austenite formed can be controlled by varying the

temperature and time of the reversion treatment as shown in Table 10.

This experiment was performed in samples measuring less than 3 mm in thickness. In larger sections the proportion of austenite formed is generally

higher due to slower cooling rate.

P29-18
19

Table 10. Influence of Reheating on the Percent Austenite

Formed in the Duplex Alloy No. 3 (ref. 19)

Reheat Temperature (°C) % Austenite

1250°C/30 mins /IBQ* 16.5**

150° 10.5

200° 14

400° 18

600° 15.5

700° 51

800° 62

1000° 68.5

All specimens were solution treated at 1250°C for 30 minutes

in silica tube (under a partial atmosphere of argon) , quenched

into iced brine and then reheated at the temperatures shown

for one hour followed by water quenching.

* IBQ: Ice Brine Quench

** Quenching out of silica tube gives a slower cooling rate

than a direct quench from air and this has caused some

austenite precipitation during cooling.

Tensile and impact data have been measured for this alloy in the hot

rolled duplex state and are given in Tables 11 and 12, respectively.

Table 11. Tensile Data at Room Temperature for Alloy No. 3

in the Hot Rolled Condition (ref. 19)

UTS (MPa) 0.2Z PROOF

STRESS (MPa)

% REDUCTION

IN AREA

Z ELONGATION

(15 um gauge length)

625 485 80 3Z

P29-19
20

Table 12. Energy Absorbed in Charpy V-Notch Test for

Alloy No. 3 in the Hot Rolled Condition

(ref. 19)

TEMPERATURE (°C)

IMPACT ENERGY

(JOULES)

20 193

0 193

-40 148

-70 102

-85 98

-196 32

These figures show acceptable strength levels coupled with fairly good

toughness down to at least -85°C (98 Joules). At -85°C, the predominant

failure mode was trans granular ductile by microvoid coalescence. Soma

small areas of cleavage were observed and presumably these are associated

with the ductile—to-brittla transition within the ferrite. It is clear,

however, that the predominantly- austenitic matrix, restricts the propagation of any- cracks which originate in the ferritic regions.

The absence of a rapid ductile-to-brittle transition in this alloy

unlike that observed in certain ferritic steels suggests that alloys of

this type could be useful in certain cryogenic applications.

Some limited impact energy data have been measured in samples containing about 25Z austenite obtained by a reversion treatment mentioned

above. The impact energy of this specimen at -40 °C was only about 27 Joules

compared to 148 Joules for tha hot rolled sample (Table 12) with 73Z

austenite. This is consistent with previous observations of decreasin g

ductility with decreasing austenite content of these alloys.

^

The

fracture mode of the sample with 25Z austenite was predominantly brittle

P29-20
21

transgranular cleavage with small areas of ductile fracture around reformed

austenite forming a continuous network along the prior ferrite grain

boundaries. This amount of austenite is apparently not sufficient to arrest

brittle cracks originating from the ferritic matrix. This result again

emphasizes the inherently brittle nature of the ferritic Fe-Mn-Al alloys.

The Austenitic Alloys :

A large proportion of this work has been centered on austenitic alloys

with compositions close to Fe-30Mn-8Al-lC . Transmission electron microscopy

was used to help characterize the aging behavior of these alloys.

Hardness versus Aging Time curves were produced for Alloys 2 and 5 at

temperatures between 450° and 650°C. In the case of Alloy No. 2, Figure 17,

the classic type of age-hardening response was observed with lower temperatures eventually achieving higher hardness values but at greatly increased

times.

For Alloy No. 5, Figure 18, aging between 450 °C and 550 °C showed the

standard type of hardness-time curves, while above 600°C, the maximum

hardness was lower and occurred after a longer time due to overaging.

In Alloy No. 5, the hardness values were always somewhat lower than

those measured for Alloy No. 2 at the same temperature. The solution

treatment temperature in both cases was 1050°C for one hour followed by

water quenching. It is not yet clear why there is a difference in the

aging behavior of the two alloys; the only compositional difference being

the carbon content.

The transmission electron micrographs in Figures 19-23 show the fine

structure which develops in this material on aging. Selected area

diffraction analysis has indicated that these precipitates form along <100>

directions in the austenite matrix and exhibit a cube-cube orientation with

P29-21
22

the austenite, Figure 20. Selected area diffraction analysis also shows the

presence of super-lattice spots and that the precipitate has a simple cubic

structure and is likely to be ordered. This has been confirmed by x-ray

analysis of the extracted precipitate particles, from which a lattice

parameter of 379.6 pm was measured and this is in good agreement with the

published data for Iron Aluminum Carbide Fe3AlC, although the interplanar

spacings for Manganese Aluminum Carbide Mn^AlC are also similar. It would

appear that the precipitate is a Fe3AlC type in which there has been some

partial substitution of manganese for iron. This has been further confirmed

by energy dispersive x-ray analysis (EDX) of the precipitate powder in the

scanning electron microscope where iron, manganese and aluminum peaks were

detected. (EDX analysis does not detect elements below sodium and therefore

carbon cannot be observed by this method.

)

The tensile and impact properties of the two alloys both in the solution

treated (ST) and quenched condition and after aging at close to the peak

hardness values have been measured. Table 13, and are similar to those

determined by other workers.

The effect of aging time at 550 °C (Alloy No. 5 after solution treatment: 1050°C/lh/WQ) on the tensile properties is shown in Figure 24. With

short aging times, it is possible to further improve the strength without

too great a loss in ductility. However, the low temperature impact

strength drops off rather drastically on aging.

For example, the room temperature impact energy after 1 hour aging

at 550°C is still a respectable 119 Joules (Table 13) at -196°C it drops

to 5 Joules. The impact energy of the solution treated material shows a

small drop from 206 Joules at room temperature to 192 Joules at -90°C and

then falls to 106 Joules at -196°C — still a very high value and a low

P29-22
Table

13.

Tensile

and

Impact

Data

for

Alloy

No. 5

and

No. 2

(ref.

19) §s

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P29-23
24

ductile-brittle transition temperature. Therefore, the increases in strength

which occur with aging are offset by a substantial decrease in toughness

together with a measurable loss of ductility as well. The aged material

failed in a predominantly intergranular brittle mode and was caused by the

grain boundary precipitates which are present in the peak-hardened state.

2. University of Calgary, Canada - W. J. D. Shaw et. al. (ref. 20)

The work has just started this summer focusing on the creep behavior

of an Fe-Mn-Al steel. The high temperature environmental (in steam) creep

and electrochemical corrosion behavior will also be studied. Preliminary

results of the short time creep tests are shown in Table 14 for an alloy

(#SB-8-24-2) containing 30.1Mn-8.3Al-1.5Si-1.07C-Fe bal. in the hot rolled

condition. These early results indicate the creep characteristics of this

alloy to be equivalent to 9Cr-lMo steels but infoerior to 304 SS. The microstructural characterization and the corrosion tests will be starting shortly.

3. Virginia Polytechnic Institute and State University, Blacksburg, VA -

J. H. Wilson, T. Sudarshan et. al. (ref. 21)

This work is also just getting underway. The major emphasis here will

be to study the effects of humidity and hydrogen on the tensile strength

and torsional fatigue life of Al-Mn stainless steels. In addition, the

effects of oxygen and hydrogen on the crack growth rates will also be

determined. The initial alloys being tested have the following compositions

(Table 15).

Table 15. Composition of Fe-Mn-Al Alloys

Being Studied at VPI&SU

Alloy %Mn %A1 %Si %C %Fe

SB-8-24-1 30.1 8.3 0.12 1.07 bal.

SB-8-24-2 30.1 8.3 1.54 1.07 bal.

P29-24
Sept.

15/82

Table

14

Dr.

W.J.D.

Shaw

Dept,

of

Much.

Engg.

Mn-Al

Steel

Univ.

of

Calgary

SB-8-24-2

Preliminary

Creep

Results

25

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P29-25

Stopped

prior

to

failure

occurring
26

The first samples cathodically charged with hydrogen and also charged

during tensile straining showed approximately the same degree of susceptibility to hydrogen embrittlement as most common austenitic stainless steels.

Intergranular fracture, typical of hydrogen-induced cracking was observed,

Figure 25. Utility of these steels in food processing and fertilizer

industries is envisioned depending on the results of this work.

4. EG&G Idaho, Inc., Idaho Falls, ID - G. Korth (ref. 22)

A major program on Mn-Al steels has also started recently at EG&G

Idaho. This will be a rather comprehensive study involving tensile, impact,

creep, fatigue, corrosion, stress-corrosion and weldability properties of

these steels. Two alloys with about 30Mn/8Al/lC/Fe, one without Si and

one with 1.5% Si have been produced for the initial tests. Early indication of weldability is encouraging. Some corrosion test coupons in geothermal wells are being tested. Other experimental work is just beginning.

SUMMARY

In summary, it is gratifying to note that considerable progress has

been made over the past year and a half on the evaluation and characterization of Fe-Mn-Al alloys. However, most of the work done and the information

presented in this report are somewhat preliminary in nature. Much work

still remains to be done for complete evaluation and understanding of various

properties of this alloy system. It appears that in certain applications

such as heat resistant materials at moderate temperatures, corrosion resistance in some environments, higher strength-to-weight ratio materials, and

cryogenic applications, etc., these alloys offer good substitution potentials

for chromium-bearing steels.

P29-26
27

It is hoped that the interested agencies will offer assistance

and encouragement for continued and more thorough investigations of this

new family of steels, offering _one option in_ the__ search jfor conservation

and substitution technology for chromium.

ACKNOWLEDGEMENTS

The continued interest of Dr. Allen Gray in the Fe-Mn-Al alloys as

potential substitutes for Cr-bearing steels is greatly appreciated. The

author is grateful to the ASM and the AIME for providing the preprints of

the manuscripts of the two papers by R. Wang et. al. and the paper by

K. Narasimha Rao et. al. , respectively. The efforts of the individual

contributors who have privately communicated with the author and provided

the information and results of their work presented in this report are

gratefully recognized. I also wish to express my special thanks to F.

Boratto and P. Tomas for providing the reprints of the papers published in

Brazil and Australia. Finally, I wish to acknowledge the continued

support of Foote Mineral Company and its permission to present this report.

REFERENCES

1.

S. K. Banerji; "An Update on Fe-Mn-Al Steels," Proc. of the Workshop

on "Conservation and Substitution Technology for Critical Materials,"

Vanderbilt University, Nashville, TN, USA, held June 15-17 (1981).

2. S. K. Banerji; "An Austenitic Stainless Steel Without Nickel and

Chromium," Metal Progress, p. 59, April (1978).

3. S. K. Banerji; "An Austenitic Stainless Steel Without Nickel and

Chromium," Trans. Indian Inst, of Metals, Vol. 30, No. 3, p. 186,

June (1977)

.

4. J. C. Garcia, N. Rosas and R. J. Rioja; "Development of Oxidation

Resistant Fe-Mn-Al Alloys," Metal Progress, p. 47, August (1982).

5. R. Wang and F. H. Beck; "An Austenitic Mn-Al-Si Steel for Propellers

of Sea-Going Vessels," to be published in Metal Progress.

P29-27
28

6. R. Wang and R. A. Rapp; "Marine Corrosion Behavior of Several Fe-Mn-Al

Austenitic Steels," to be submitted for publication.

7. K. Narasimha Rao, R. Sivakumar and M. L. Bhatia; "An Evaluation of

Al-Mn-C Austenitic Steel," submitted to Journal of Metals.

8. L. C. Casteletti and D. Spinelli; "The Resistance to Oxidation and the

Electrical Resistivity of an Fe-Mn-Al System Alloy," Dept, of Matls.

Sci., EESC, Univ. of Sao Paulo, Sao Carlos, Brazil; Proc. XXXVI,

Brazilian Soc. for Metals (ABM) Annual Congress, p. 249, held in

Recife, Pernambuco, Brazil, July 5-10 (1981).

9. F. C. R. Assuncao and M. F. da S. Lopes; "Study of Oxidation of

Fe-Al-Mn-C Alloy at High Temperatures," Dept. Matls. Sci. & Met.,

Catholic Univ., Rio de Janeiro, Brazil, ibid, p. 329.

10. A. J. A. Bushinelli, J. C. Dutra and W. May; "Study of the Microstructural Transformations in TIG Welding of an Fe-Al-Mn Alloy,"

Fed. Univ. of Santa Catarina, Florianapolis, Brazil, ibid, p. 521.

11. L. C. Casteletti and D. Spinelli; "Mechanical Properties of an Austenitic Steel of the System Fe-Mn-Al," Dept. Matls. Sci., EESC,

Univ. Sao Paulo, Sao Carlos, Brazil, ibid, p. 299.

12. A. P. Tschiptschin and H. Goldenstein; "Precipitation Reactions in

Low-Carbon Fe-Mn-Al Alloys," Dept. Met. Eng., Univ. of Sao Paulo,

Brazil, ibid, p. 117.

13. J. R. T. Branco and F. J. M. Boratto; "Domain of the Austenitic Phase

in the Fe-Mn-Al System at 1000°C," Met. Tech. Section, Tech. Center

Found, of Minas Gerais-CETEC , Belo Horizonte, Brazil, ibid, p. 175.

14. P. Tomas; "Elements Essential to High Temperature Sulphidation

Resistant Iron-Based Alloys," Univ. of New South Wales, Kensington,

NSW, Australia, Proc. 35th Annual Conf. of the Australasian Inst,

of Metals, p. 90, held in Sydney, Australia, May (1982).

15. P. R. S. Jackson and G. R. Wallwork: "The Development of Fe-Al Alloys

for High Temperature Oxidation Resistance," Univ. of New South Wales,

Kensington, NSW, Australia, ibid, p. 78.

16. D. J. Chakrabarti; "Phase Stability in Ternary Systems of Transition

Elements with Aluminum," Met. Trans. B, Vol. 8B, p. 121, March (1977).

17. P. Tomas and S. K. Banerji, private communication (1982).

18. R. A. Perkins; "Sulphidation-Resistant Alloy for Coal Gasification

Service," Qtrly. Report #FE-2200-6, Lockheed, Palo Alto, CA (1976).

19. G. E. Hale, Dept, of Metallurgy, University of Leeds, Leeds, U.K.;

Ph.D. thesis to be submitted, private communication (1982).

P29-28
29

20. W. J. D. Shaw, Dept, of Mech. Eng., University of Calgary, Calgary,

Alberta, Canada; private communication (1982)

.

21. T. Sudarshan and J. H. Wilson, Depts. of Materials and Agricultural

Eng., VPI&SU, Blacksburg, VA, private communication (1982).

22. G. E. Korth, EG&G Idaho, Inc., Idaho Falls, ID, private communication

(1982).

SKBanerjiierd

10/82

P29-29
30

Weight gain, g/m*

Fig. 1. Plots of weight gained versus time at 700°C (1290°F) for the six

experimental alloys of Table 1, type 304 stainless steel, and a 1010 carbon steel. Parabolic regions of the data curves are presented here. Note that the slope for alloy B is similar to

that for type 304. Pressure: 1 atm (100 kPa) of flowing oxygen,

(ref. 4)

P29-30
31

Weight gain, g/m*

Fig. 2, Plot of weight gained versus time at 500°C (930°F) for the six

experimental alloys of Table lj type 304 stainless steel, and a 1010 carbon steel. Parabolic regions of the data curves are presented here. Note excellent behavior of alloy F. Pressure:

1 atm (100 kPa) of flowing oxygen, (ref. 4)

P29-31
Fig. 4. Weight gain as a function of time at 800°C in air. (ref. 8)

P29-32
Resistivity

(jjA-cm)

33

Fig. 5. Temperature dependence of resistivity for an Fe-31Mn, 7.5A1,

1.3Si, 0.93C alloy, (ref. 8)

P29-33
34

Fig. 6. Variation in tensile strength with temperature, (ref. 11)

P29-34
Yield

strength

(MN/m*)

35

Fig. 7. Variation in yield strength with temperature, (ref. 11)

P29-35
36

Fig. 8. Variation in tensile strength with the degree of reduction,

(ref. 11)

P29-36
37

Fig. 9. Variation in yield strength with the degree of reduction,

(ref. 11)

P29-37
Elongation

(%)

38

Fig. 10. Variation in elongation with the degree of reduction,

(ref. 11)

P29-38
Hardness

fe.H )

39

Fig. 11. Variation in hardness with the degree of reduction, (ref. 11)

P29-39
Energy

absorbed

in

fracture

(J)

40

Fig. 12. Variation in the energy absorbed in impact with temperature,

(ref. U)

P29-40
41

WEIGHT GAIN mgm.cm.*2

Fig. 13. A compilation of the weight gain data after 9 hours reaction in a gas mixture of hydrogen plus 1 volume % hydrogen sulphide;

pS 2 = 3 x 10“ 9 Atm at 700°C and pS£ = 2.4 x 10”8 Atm at 800°C.

(ref. 14, 17)

P29-41
42

Matrix

Fig. 14. Duplex Scale formation on Femnal 30 after 9 hours

in H 2 /I vol. % H 2 S at 700°C. Cross section of

scale showing duplex layers. The outer layer is

predominantly MnS and the peg-like inner layer is

aluminum sulphide, (ref. 14)

P29-42
43

Fig. 15. A comparison of the scale topographies present on two

different Fe-Al-Mn-C alloys, (a) Femnal 30, and

(b) 11-10-4B; 9 hours at 800°C in H£ - 1 vol. % I^S.

(ref. 17)

P29-43
44

Fig. 16. Ferritic Alloy No. 4 (9Mn/7Al/ .07C/Fe) - as cast state, (ref. 19)

P29-44
Fig.

17.

AGEING

CURVES

FOR

AN

Fe-30.4Mn-7.6Al-0.81C

ALLOY

(No.

2)

AFTER

SOLUTION

TREATMENT

AT

1050°C

FOR

ONE

HOUR

FOLLOWED

BY

WATER

OUENCHTNn

. (

ref

#

lg)

45

P29-45

AGEING

TIME!

MINUTES)
Fig.

18.

AGEING

CURVES

FOR

AN

Fe-30.9Mn-8.0Al-0.97C

ALLOY

(No.

5)

AFTER

SOLUTION

TREATMENT

AT

1050°C

FOR

ONE

HOUR

FOLLOWED

BY

WATER

QUENCHING

.

(ref.

19)

46

I lf

0Oo o c_> u o

o °o °o °o °o

in o in o in

"j’ m in <JD to

02 AH SS3IM0UVH

P29-46
47

Fig. 19. Alloy No. 2.

Hot rolled bar - ST 950°C/30 mins/water quench (WQ)

Cold rolled 90%

ST 950°C/1 hr/WQ

Age 600° C/3 mins/WQ

Note the presence of very fine precipitation after

aging for only three minutes, (ref. 19)

P29-47
48

Fig. 20. Alloy No. 2. Treatment as in Fig. 19 but aged at 600°C

for 30 minutes and water quenched.

Dark field micrograph from (010) precipitate spot in

Fig. 21. (ref. 19)

Fig. 21. Alloy No. 2. Selected area diffraction pattern from area seen

in Fig. 20 showing superlattice spots intermediate between the

austenite matrix spots, (ref. 19)

P29-48
49

Fig. 22. Alloy No. 2. Treatment as in Fig. 19 but then aged at 600°C

for one hour and water quenched. Bright Field Micrograph,

(ref. 19)

Fig. 23. Alloy No. 2. Dark Field Micrograph of area seen in Fig. 22.

Zone axis just of [001]y. Operative reflection is a (100)

precipitate spot. (ref. 19)

P29-49
and

0.2%

Proof

Stress

(MPa)

50

5

1100 .

1000

900

800

700

600

500

400

0

ST+Q

10 20 30 40

Aging Time (Hours)

50

80

60

40

20

0

Fig. 24. The effect of aging time (at 550°C) on the tensile properties of

Alloy No. 5 after ST 1050°C/lHr/WQ . (ref. 19)

P29-50

%

Reduction

in

Area

and

Elongation
51

VRJ&SU C

Fig. 25. Hydrogen-induced intergranular cracking in an Fe-Mn-Al alloy, (ref. 21)

P29-51

THE DEVELOPMENT OF NEW ALLOYS

TO REPLACE CHROMIUM IN CARBURIZING STEELS FOR GEARS AND SHAFTS

CARL J. KEITH AND V. K. SHARMA

Materials Engineering and Technology

International Harvester Company

Hinsdale, Illinois

ABSTRACT

A status report on the Bureau of Mines sponsored research to develop a new series of cost-effective chromium-free steels is provided. These

new steels will be capable of adoption by a broad range of industry as a

substitute for the chromium-containing 8600 and 4100 standard grades

of heat treated constructional alloy steels. Based upon several socioeconomic scenarios and their projected impact on alloying element cost

and availability in 1985, International Harvester's computerized

metallurgical design system has been used to design two substitute

steels. Both steels, a manganese-molybdenum substitute and a manganese-nickel-molybdenum substitute, are expected to provide

microstructure, heat treat response, and mechanical properties

equivalent to the 8600 and 4100 steels. Small experimental heats of

both chromium-free replacement steels were prepared and their

hardenability response and mechanical properties were evaluated and

compared to the base line 8620 steel.

P30-1
INTRODUCTION

Chromium is both a critical and strategic alloy element. With the exception of

approximately 8 percent chromium recycled from stainless steel scrap, the U.S. relies

totally on imports. Chromium resources in the Western Hemisphere are limited to about three-tenth of one percent of the world resources of chromite ore. Identified

world supplies of chromium are sufficient to meet conceivable demand for many

centuries, however 97 percent of the world's chromite reserves and 98 percent of its identified resources are located in Rhodesia (Zimbabwe) and South Africa^). This very

high concentration of world reserves in a politically unstable area, coupled with the

relatively rapid rise in world demand and limited substitution potential, creates a potentially dangerous situation. The National Materials Advisory Board in their report on contingency plans for chromium utilization points out that "at some future time the

major source of chromite could become inaccessible to the United States as a result of

political action."(2)

Depending on the level of economic activity, sixty-to-seventy percent of chromium

consumed annually in the United States is used for metallurgical purposes. The bulk of

this usage is in the production of heat resistant steels, superalloys, and corrosionresistance stainless steels. At the present time there exists no viable substitute for

chromium as used in these heat and corrosion resistant applications. However,

approximately 17 percent of the chromium consumed annually is used in the production

of the constructional alloy steels. In these applications chromium is used primarily for

its effect on hardenability and, while its use is both highly efficient and cost-effective,

it can be substituted for. Two particular grades of constructional alloy steel,

chromium-molybdenum (4100) and nickel-chromium-molybdenum (8600), account for

60 percent of the chromium used in the constructional alloy steels and thus

approximately 10 percent of the total U.S. demand. Conservation by substitution of

this chromium could reduce demand by 10 percent or alternatively expand existing

supply thereby extending the supply for those heat- and corrosion-resistant applications

for which there is currently no viable substitute for chromium.

The purpose of the Bureau of Mines sponsored research described in this paper is to

develop a new series of chromium-free constructional alloy steels capable of replacing

the two chromium-containing AISI 8600 and 4100 steels. The goal is to substitute for

the chromium in these two steels other cost-effective alloying elements in which the

U.S. is self-sufficient or for which a more secure and stable source of supply exists. In the event of a chrome supply disruption or significant price escalation in the time

period 1985-1995, the new chrome-free compositions can be adopted by a broad range

of U.S. industry without significant testing and development effort. To be readily

acceptable to domestic steel producers, the new chrome-free steels have been designed

to avoid the necessity for special production practices which are not currently required

for the production of the standard grades. For example, although it is more

cost-effective to design chrome-free steels with manganese level exceeding 1.30

percent, the manganese range was limited to 1.00 - 1.30 percent because a broad range

of domestic steel producers are not accustomed to producing higher manganese steel on

a commercial basis.

P30-2
SUBSTITUTE STEELS

The acceptance and adoption of new steels as standard grades is a relatively lengthy

process requiring many years. Most of the currently used constructional alloy steel

grades were developed in the late 1930s and adopted in the early 1940s as National

Emergency Steels during World War II. Engineering design of parts and components,

using these steels is well advanced, and manufacturing methods rely on process control

parameters developed only after long and painful experience. It would be very difficult

to replace parts made from these standard steels with parts made from new alloy

compositions having different hardenability, heat treat response, and mechanical

properties without redesigning the parts. Redesign of the parts would involve extensive

engineering and manufacturing test programs. Engineering test programs would be

required to ensure achievement of required engineering performance, and reliability and

manufacturing test programs would be required to establish new process control

parameters to enable consistent production of parts within required tolerance levels.

However, experience has shown that standard steels can be replaced by new steels of

different alloy composition provided that the new steels exhibit the same hardenability,

heat-treat response and mechanical properties. By careful control and balance of the

composition, new steels can be developed to replace the standard steels in current parts

without the need for redesign of the parts and with only nominal changes in process

control parameters. Basic assumption in the development of such substitute steel is that the engineering performance of a heat treated component is controlled by the

carbon content, microstructure, and residual stresses. The microstructure and residual

stresses are governed by the carbon content, hardenability, and martensitic start-andfinish transformation temperatures. Alloy composition controls hardenability and

transformation temperatures. Opportunities, therefore, exist to develop substitute

steels composed of alternate cost-effective and/or readily available non-strategic alloy

elements provided the new steels have the same carbon content, hardenability, and

martensitic transformation temperatures.

The use of substitute steels as replacements for the standard steel grades is not a new

concept. The SAE specification 31081 was developed to provide a uniform means of

designating new steels during a period of usage prior to their acceptance as standard

steel grades. Beginning in the early 1960s, new replacement steel grades developed by

industry have been listed in this specification as EX steel grades. At International

Harvester a computerized metallurgical design system has been used successfully over

the past ten years to develop a number of proprietary, minimal cost, alloy compositions

which are presently widely used in the production of IH products. None of these new

steels is a chromium-free steel. At the present time, the use of chromium to enhance

hardenability is highly cost-effective and the wide spread implementation of chromium

substitution would not be in the economic best interest of the nation. As noted earlier,

in the design of new chromium-free steels alloy costs have been minimized within the

constraints of self-sufficiency and/or availability considerations.

Computer-harmonizing or the CH portion of International Harvester's metallurgical

alloy steel design system CHAT (Computer Harmonized Application Tailored) was used

to design new chromium-free steels. The IH CHAT system has been fully described

elsewhere (3,4). The CH portion of the system is a computer program which uses linear

and separable programming to optimize a primary or objective function while at the

same time satisfying a series of additional functions. When the CH program is used to

design replacement steel grades, the primary or objective function is described as the

total cost of the alloy additives required to make the steel. This primary function is

P30-3
minimized, while at the same time a series of restrictive functions embodying minimum

and/or maximum base and case hardenability requirements, transformation temperature

requirements, and additional limitations aimed at controlling intergranular oxide

penetration during heat treatment and carbide formation during carburizing are all simultaneously satisfied.

ALLOY STEEL COST FACTORS

One of the key pieces of information required in designing a cost-effective replacement

steel is the alloy steel cost factors. The cost factors used in the CH portion of the

CHAT system are the raw material costs to the steel mill to add one percent of the

alloy element to a full heat of steel. These cost factors are modified for the size of the

heat and, the yield and recovery factors which account for oxidation, metal slag

reactions, etc. and affect the efficiency of the alloy element additions. The CHAT

system selects the optimized combination of alloy elements (minimized cost) meeting

the input requirements and limitations. The resulting CHAT composition is then

modified to provide an optimum price in terms of its "chemistry grade extra" as determined by commercial steel product price books. This difference between steel

cost and steel price must be clearly appreciated. Although there is a close correlation

between steel mill cost and steel product price, there is not always a one-to-one

correspondence.

In designing the chromium-free steels, the 1981 alloy cost factors were projected to

1985 and then modified to reflect the following four future chromium supply scenarios

considered probable in 1985:

Scenario A: Chromium supply will remain stable with prices rising in a steady manner

as world economic and political conditions continue to evolve in a relatively orderly

fashion. Based upon world conditions existing at this time, this scenario appears to be

the most likely and is estimated to have a relatively high 75 to 80 percent probability.

Scenario B; Chromium remains available but supply and price fluctuate as the

principal African producers seek to form an OPEC-like cartel and attempt to control

pricing. This scenario is considered to be eminently possible due to the concentration

of resources in southern Africa. However, this scenario is assigned a relatively low

probability of 15% because of the high dependency of that area on export of minerals

for foreign trade; the relatively low economic development of the area; and the

reasonable availability of lower grade ores from other sources (including industrial and

governmental stock piles) for both alloy element and ferro alloys.

Scenario C: Chromium supply from African sources is cut off as a result of a regional

war or other occurrences causing a complete breakdown of export capability from the

South African sources. As a result of political unrest, the potential for insurrection

and/or war in this area is very real. The relative isolation of South Africa in the United

Nations and world community and the active support of Soviet Bloc nations for the

South African opposition groups are destabilizing factors. However, the interest of the

major western powers in the mineral resources of the area and the relative economic

and military strength of the nation of South Africa in relation to its own neighbors

indicate that a war could probably be contained without seriously affecting export. The

probability of a complete breakdown of South African production and export potential is therefore considered to be relatively low, i.e. 5 percent.

P30-4
Scenario D: This scenario assumes the breakdown of African chromium supply as in

Scenario C above, but then further assumes that at the same time Soviet and Albanian

sources of chrome supply would also be cut off. This fourth scenario D is assigned a very low probability of 1 percent.

The costs for the alloy elements developed from various scenarios are shown in Table I. The cost forecasts, developed by the Resources and Technology Forecasting Group of

IHC Corporate Technology as a part of an ongoing study, are intended to represent

long-term trends rather than short-range fluctuations.

TABLE I - STEEL ALLOY COST FORECAST FOR 1985

Unit = $/ton (2,000 lbs)

Ni* Mo* Cr* Mn* Ferroalloys

Scenario NiO

do

M0O2 Cr2C>3 MnO FeCr FeMn FeSi

Scenario A 9,760 31,240 489 97.2 1,431 750 1,240.

Scenario B 10,736 34,364 562 107.0 1,646 862 1,302

Scenario C 11,712 37,488 3,600 116.6 3,856 937 1,364

Scenario D 11,712 37,488 3,960 116.6 4,242 937 1,364

1981 Prices

(Nov-Dec 1981) 7,160 21,340 360 71.5 1,067 538 900

Price per ton of Contained Metal:

Scenario A: Stable world conditions, 75-80% probability

Scenario B: Producers' Cartel, 15% probability

Scenario C: South African Breakdown, 5% probability

Scenario D: Scenario C plus Soviet and Albanian sources cut off, probability 1%

The alloy element cost forecasts were reviewed with a number of experts in various areas of alloy availability. Of particular interest were comments from Climax

Molybdenum Corporation regarding the projected costs for molybdenum oxide. The

forecasts, conducted at a time when molybdenum was still considered to be in short

supply, projected a 1985 molybdenum price tending toward $ 1 6/lb under stable world

conditions of Scenario A. The Climax Molybdenum representatives suggested that the

existing 1981 producer price of $6. 85/lb would prevail throughout 1982 and would not

exceed $ 1 0/lb by 1985. It was pointed out that the continuing economic recession,

depressed steel market demand, and increased world production had affected producer

inventory levels to a point where they could not be adjusted in the near-term future.

Figure 1 and Table II reported in American Metal Market, July 1982,^) tend to confirm

this outlook.

P30-5
AMM—MOLYBDENUM SUPPLEMENT Figure 1 - Molybdenum Price and U.S. Inventory(^)

TABLE II - THE OUTLOOK FOR MOLYBDENUM^)

(in thousands of metric tons)

1982 1985 1990

Consumption

United States 25.4 34.7 36.5

Other Countries 43.6 60.0 80.8

TOTAL 69.0 94.7 117.3

Supply (Mine)

Capacity 117.5 140.3 156.2

Operating Rate .69 .74 .84

Production 81.7 103.7 131.6

Price of Oxide ($/lb)

1980 $

$ 5.60 $ 6.80 $ 7.93

The cost factors used in the CH computer program significantly affect the selection of

the optimum composition. To more fully evaluate the problems of cost projections, a

cost sensitivity analysis for two competitive alloying elements, nickel and molybdenum,

was performed. This cost sensitivity study included the constraints imposed in

matching the metallurgical characteristic base and case hardenability of the 8600 and

P30-6
4100 series steels, and is not applicable to other steel grades having different

metallurgical characteristics. The results indicate that for the 8600 and 4100 series

steels a Mo/Ni cost ratio of 2.4 is a critical point. Above this ratio (i.e.. Mo cost

greater than 2.4 times Ni cost), nickel substitutes to reduce Mo content; below this

ratio, nickel does not substitute for Mo. A review of the history of alloy costs over the

past 20 years indicates that the Mo/Ni cost ratio has existed on both sides of the 2.4

critical ratio. As it will be seen later, it is precisely for this reason that two different

chrome-free compositions, a manganese-nickel-molybdenum and a manganese-molybdenum steel, have been developed in this program.

COMPOSITION OF CHROMIUM-FREE STEELS

Manganese, nickel and molybdenum, are three possible elements which can compensate

for the base and case hardenabilities lost due to the removal of chromium in the

chrome-free steels. Increasing the base carbon can increase the base hardenability, but

in designing substitute steel only a minimal change in the base carbon is permitted.

Obviously a change in the base carbon will have no effect on the case hardenability.

Silicon contributes to case hardenability significantly but has no significant effect on base hardenability unless increased beyond approximately 0.60%. Silicon range in

standard AISI steels is 0.15 to 0.35%. Producers and users of carburizing steels do not

have experience with higher silicon carburizing grades of steels. Therefore, three key

elements which can be effectively used to replace chromium are manganese, nickel, and

molybdenum.

Although manganese is a relatively inexpensive alloying element, consideration was

given to the use of both high and low manganese in the computer design of the

chromium-free replacement steels. South Africa currently supplies approximately 35%

of the Western demand for manganese(% and the formation of a producer cartel for

chromium which included the South African sources for chromium would in all probability affect manganese also. A breakdown of the African export capability for

chromium would also involve export capability for manganese. These factors, and their

effects on manganese costs, are included in Table I, Steel Alloy Cost Factors, under

Scanarios B, C, and D. However, it is assumed that short-term manganese shortages

could be offset by increased production from other western suppliers (primarily

Australia and Brazil).

Four different manganese ranges varying from 0.40 to 1.50 percent Mn were considered.

These four manganese ranges, along with additional restrictions of fixed silicon range, no chromium, and metallurgical base and case hardenability requirements for 8622 and

4118 steels, were included as input to the CH program. Using the 1985 alloy steel cost

factors for various scenarios, the computerized alloy steel design system provided

optimal chromium-free compositions for each manganese range. Chromium-free

replacement steel compositions developed for the AISI 8620 steel are shown in Table III. The CH program first provided only the -1/4 chemistry. The +1/4 chemistry was

established after a due consideration to the chemical range limits (tolerances) for the

various alloying elements established by the American Iron and Steel Institute (AISI).

This was accomplished within the CHAT system by varying the input requirements.

Input requirements representing the minimum, maximum, and/or median values of the

metallurgical characteristics describing the steel were varied systematically until the

resulting new composition met AISI range tolerances and the primary objective function

was optimized. Notice that the chromium-free chemistries in Table III are applicable

to all four scenarios. Changes in the alloy cost factors resulting from various scenarios

did not alter the composition. The computer selected the same cost-effective

chemistry irrespective of alloy cost factors in Scenarios A to D, which indicates that

P30-7
TABLE III COMPOSITIONS OF CHROME-FREE REPLACEMENTS FOR AISI 8620 STEEL

ITEMS

AISI 8620

CH' VERSION SELECTION - STANDARD RANGES

.40-. 60% Mn .70-.90% Mn 1.00-1.30% Mn 1.20-1.50% Mn

-1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4

C .18 .23 .18 .23 .18 .23 .18 . 23 .18 .23

Mn .75 .85 .45 .55 .75 .85 1.07 1.23 1.27 1.43

Si .20 .30 .20 .30 .20 .30 .20 .30 .20 .30

S .03 .05 - - - - - - - - Cr .45 .55 .02 .08 .02 .08 .02 .08 .02 .08

P .01 .03 - - - - - - - - Ni .47 .63 .01 .05 .35 .55 .35 .55 .25 .35

Mo .17 .23 .72 .83 .42 .48 .27 .33 .27 .33

Min 1.42 1.30 1.41 1.53 1.71

D^Max 2.33 2.06 2.19 2.45 2.64

D,

c Min 4.31 5.03 4.15 4.09 4.39

Dt max

Ic 6.45 7.08 5.83 6.21 6.23

Ms Base 797° 745° 8300 7820 807° 7550 791° 7360 7840 7310

Mfo 779° 7270 8120 7640 7890 7370 7730 7180 7660 7130

m50 712° 6610 7450 6980 7220 6700 707 0 6510 6990 6470

M90 612° 5600 6450 5970 6210 5690 6060 5510 5980 5460

Mf 410° 358° 443° 3950 4200 368° 4040 3490 3970 3440

Ms Case 305° 2870 343° 3300 3210 3040 3050 2840 2960 2780

Acl 1329© 1330° 13360 13390 13190 13180 1313° 1311° 13120 13130

ac3 1528° 1515° 15720 1565° 1545° 15310 15370 15230 15400 15280

Ae3 1493° 1477° 1516° 15040 1498° 1481° 1490° 1471° 1488° 1473°

Cost Factor

$ - 1981 $9.50/cwt $19 .50/cwt $14.85/cwt $10.85/cwt $10 .05/cwt

$

- 1985A $ 12 . 10 /cwt - - $13.20/cwt $12.60/cwt

$ - 1985B $14.90/cwt - - $15.25/cwt $13.85/cwt

$ - 1985C $17.80/cwt - - $15.90/cwt $15.30/cwt

$

- 1985D $22.40/cwt - - $15.90/cwt $15.30/cwt

Note: Df values calculated from +1/4 chemistry.

P30-8
under the constraints imposed relative per dollar contributions of various alloying

elements to hardenability does not change with the increase in the alloy cost factor

from Scenario A to D.

Table III shows +1/4 chemistries, characteristic hardenabilities, and various transformation temperature for the 8620 and the computer harmonized replacement versions for

the four manganese ranges. Present and 1985 estimated prices under the four scenarios are given in the last column. The .40 - .60 Mn range results in the development of a Mn-Mo steel with a .70 -.85 percent molybdenum range. It would be a costly steel, and

its use would require more extensive testing. The .70 - .90 Mn range results in the

development of a Mn-Ni-Mo steel, which again is relatively costly. The 1.20 - 1.50 Mn

range results in the development of the most cost-effective compositions. However, in

current steel production practice, the use of manganese above 1.20 to 1.30 percent

involves additional problems in ingot casting practice due to a tendency of the high

manganese steel to cause cracking if allowed to cool in the ingot. This problem can be

avoided by vacuum degassing and/or by controlled processing in the steel mill. Since

the chromium-free steel to be developed in this project is intended for use by a broad

range of industry, without a significant change in current production practices, the

1.20 - 1.50 percent Mn range is rejected, as are the .40 - .60 percent Mn and .70 - .90 percent Mn ranges, because they are too expensive. Based on these reasons, the

1.00 -1.30 percent Mn range which results in a replacement steel slightly more

expensive than the 1.20 - 1.50 percent Mn steel, is selected as the optimal manganese

range.

The 1.00 - 1.30 percent Mn steel CH composition was further refined to provide an optimum price in terms of the chemistry grade extras and to reproduce all metallurgical requirements of hardenability and transformation temperatures of the parent 8620

steel as closely as possible. Table IV shows the revised Mn-Ni-Mo composition. In Table III, for the 1.00 -1.30 percent manganese steel, the computer indicates an optimal

cost with a steel containing nickel at a range of .35 - .55 percent. However,

commercial practice does not distinguish (provides no price break) a difference on

nickel content below a value of 0.7 percent. Therefore the nickel content in Table IV

was increased to a range of .55 - .65 percent without affecting cost. This increase

enhances case hardenability without significant increase in the base hardenability.

Another modification in the chemistry is a decrease in the carbon range from 0.18 - 0.23 percent to 0.16 - .21 percent. Examination of the output of the CH program,

Table III, indicates that the high carbon, case hardenability (Dic) is the primary factor

influencing the alloy composition. In order to meet the minimum Dic requirement, the

base hardenability, D15 , is being exceeded. Experience has shown that although the

increased base hardenability does not affect the engineering performance of the new

steel, manufacturing process problems are sometimes encountered as a result of

increased core hardness in comparison to the previously used standard steel. The

problem of increased core hardness or base hardenability can be ameliorated by a slight

reduction in the carbon content. This reduction in the carbon content reduces the base

hardenability without affecting the case hardenability. The resulting calculated 3ominy

hardenability band for the chrome-free steel matches the standard 8620 very closely, as shown in Figure 2.

P30-9
TABLE IV

COMPOSITIONS OF CHROME-FREE REPLACEMENTS FOR AISI 8620 STEEL

Items

AISI 8620

CH" - REVISIONS

1.00-1.30% Mn

Mn-Ni-Mo

1.00-1.30% Mn

Mn-Mo

-1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4

C .18 .23 .16 .21 .16 .21

Mn .75 .85 1.07 1.23 1.07 1.23

Si .20 .30 .20 .30 .20 .30

S .03 .05 - - - - Cr .45 .55 .02 .08 .02 .08

P .01 .03 - - - - Ni .47 .63 .55 .65 .01 .05

Mo .17 .23 .27 .33 .37 .43

Djb Min 1.42 1.55 1.52

Djb Max 2.33 2.28 2.19

Djc Min 4.31 4.42 4.43

Djc Max 6.45 6.41 6.11

Ms Base 797° 745° 800° 7 48° 816® 7660

M 10 779° 727° 7820 7300 7980 7480

m50 712° 661° 7160 663° 7310 6810

M90 612° 5600 6150 5630 6310 5800

Mf 4 1

0

o 358° 413° 361° 429° 3790

Ms Case 305° 287° 299° 2990 313° 297°

^cl 13290 1330° 1307° 1307° 1324° 1326°

^c3 1528° 1510O 1540° 1540° 1561° 1550°

Ae3 1493° 1477° 1490° 1490° 1507° 1494°

Cost Factor

$ - 1981 $9.50/cwt $10.85/cwt $9.60/cwt

$ - 1985A $ 12 . 10 /cwt $13. 20/cwt $1 1 .85/cwt

$ - 1985B $14.90/cwt $15.25/cwt $12.95/cwt

$ - 1985C $17.80/cwt $15.90/cwt $14.30/cwt

$ - $1985D $22.40/cwt $15.90/cwt $14.30/cwt

NOTE: Dj values calculated from +1/4 chemistry.

P30-10
(

n

in

tUzQ

cc<

I

o

til §

o

o

cr

8

C

3 2 0 H UCU1

-BAN

A T E D

3

H - B AND ... Ov

V

\

<•X A

\

\l

'i

v»\\

\xl

k>

•’H1 *

“’ —' —— - —

4 8 12 16 20 24 28 32

JOMINY DISTANCE , 1/16 INCH

Figure 2 - Standard 3ominy hardenability band for AISI 8620

steel compared to the calculated hardenability band for the

chrome-free replacement steels.

It should be pointed out that the Mo/Ni price ratio used in this portion of the CH

development program is approximately 3.2, as shown in Table I. Consideration of a Mo/Ni price ratio below 2.4, which actually exists at the time of writing this report,

results in the development of the manganese-molybdenum composition shown in

Table IV as the most cost-effective replacement steel. Since the Mo/Ni price ratio has

historically varied on both sides of the critical ratio of 2.4, both the Mn-Ni-Mo and the

Mn-Mo chromium-free steels are considered as future replacement compositions. The

selection of the most cost-effective replacement will depend upon the relative costs of

nickel and molybdenum at that future time when the need for the chromium-free steel occurs. The calculated Jominy hardenability band for the Mn-Ni-Mo and for the Mn-Mo

steel is represented by the same curves in Figure 2.

Similar to the 8620 steel, CH development of chromium-free replacements for a representative 4100 series steel (4118) is given in Tables V and VI. The ladle chemical

analysis ranges for the chromium-free replacement steels are summarized in Table VII. The calculated hardenability band for the Mn-Ni-Mo and Mn-Mo chromium-free

replacement steels are compared to the standard AISI 4118 steel hardenability band in

Figure 3.

In the development of the replacement steels, it is assumed that the chromium-free

steels will be produced using essentially BOF practice and that the steels will be ingot

cast rather than strand cast. If electric steel practice is used, the melt composition can be modified to account for the high residual element content normally present in

the scrap charge. For the steel users, the chromium-free steels will have to be

provided in a fully-killed fine grain condition with a standard silicon range, since the

8600 and 4100 steels used domestically in carburizing applications are generally

specified to be fully-killed fine grain steels.

P30-11
TABLE V

COMPOSITIONS OF CHROME-FREE REPLACEMENTS FOR AISI 4118 STEEL

ITEMS

AISI 4118

CH' VERSION SELECTION - STANDARD RANGES

.40-.6096 Mn .70-.9096 Mn • 1.00-1.30% Mn 1.20-1.50% Mn

-1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4 -1/4 + 1/4

C .18 .23 .18 .23 .18 .23 .18 .23 .18 .23

Mn .75 .85 .45 .55 .75 .85 1.07 1.23 1.27 1.43

Si .20 .30 .20 .30 .20 .30 .20 .30 .20 .30

S .03 .05 | - - - - - - - Cr .45 .55 .02 .08 .02 .08 .02 .08 .02 .08

P .01 .03 - - - - - - - Ni .01 .05 .01 .05 .35 .55 .01 .05 .01 .05

Mo .10 .13 .52 .63 .27 .33 .17 .23 .15 .18

Dj, Min 1.15 1.05 1.17 1.22 1.37

D., Max 1.76 1.68 1.83 1.87 1.99

D Min 3.17 3.75 3.18 3.02 3.30

Dj^Max 4.29 5.25 4.58 4.27 4.19

Ms Base 813° 765° 833° 785° 8090 7570 804° 753° 7930 7430

MlO 7950 747° 8150 767° 7910 7390 7860 7350 7750 7250

M50 728° 681° 748° 701° 7240 6720 7190 6690 708° 658°

M90 627° 5800 647° 6000 623° 5710 6180 5680 6070 5580

Mf 426° 378° 446° 398° 4220 370° 417° 3660 4060 3560

Ms Case 3210 3070 3470 3340 3240 3070 3170 3010 3060 2900

^cl 1343° 1348° 1336° 1339° 1319° 13180 13240 13260 13200 13220

^c3 1536° 1525° 1560° 15530 15370 15230- 15400 15310 15390 15280

^e3 1507° 1495° 1516° 1504° 1498° 1481° 1501O 1487° 1496° 14820

Cost Factor

$ - 1981 $4.90/cwt $ 12 . 10 /cwt $10.75/cwt $6 . 10 /cwt $5 .65/cwt

$

- 1985A $5.95/cwt - - $8 . 00 /cwt $7.00/cwt

$ - 1985B $7.65/cwt - - $8.95/cwt $7.70/cwt

$ - 1985C $10.50/cwt - - $9.90/cwt $8.45/cwt

$

- 1985D $15.10/cwt - - $9.90/cwt $8.45/cwt

NOTE: Dj values calculated from +1/4 chemistry.

P30-12
TABLE VI

COMPOSITIONS OF CHROME-FREE REPLACEMENTS FOR AISI 4118 STEEL

ITEMS

AISI 4

1 1 Q

CH" - REVISIONS

118

1.00-1.30% Mn

Mn-Ni-Mo

1.00-1.30% Mn

Mn-Mo

-1/4 + 1/4 -1/4 +1/4 -1/4 + 1/4

C .18 .23 .16 .21 .16 .21

Mn .75 .85 1.07 1.23 1.07 1.23

Si .20 .30 .20 .30 .20 .30

S .03 .05 - - - - Cr .45 .55 .02 .08 .02 .08

P .01 .03 - - - - Ni .01 .05 .25 .35 .01 .05

Mo .10 .13 .17 .23 .22 .28

Dib Min 1.15 1.26 1.25

Dib Max 1.76 1.85 1.81

Djc Min 3.17 3.17 3.33

Dic Max 4.29 4.74 4.70

Ms Base 8130 7650 8110 7590 8180 768°

M 10 795° 7470 7930 7410 8000 750®

m50 7280 6810 7270 6740 7340 6830

M90 627° 580° 6260 5740 633° 5820

Mf 426® 378® 4240 3720 4310 3810

MsCase 3210 307° 3100 2920 316° 300°

^cl 1343° 1348° 1310O 1317° 1324° 13260

^c3 1536° 15250 15430 15300 15520 15410

^e3 1507° 1495° 1500O 1484° 1507° 1494°

Cost Factor

$ - 1981 $4.90/cwt $7.50/cwt $7.10/cwt

$ - 1985A $5.95/cwt $9.30/cwt $8.85/cwt

$ - 1985B $7.65/cwt $10.25/cwt $9.80/cwt

$ - 1985C $10.50/cwt $1 1.30/cwt $10.65/cwt

$ - 1985D $15.10/cwt $1 1.30/cwt $10.65/cwt

NOTE: Di values calculated from +1/4 chemistry.

P30-13
TABLE VII CHROME-FREE REPLACEMENT COMPOSITIONS FOR STANDARD 4118 AND 8620 STEELS

LADLE ANALYSIS RANGES

Chemistry

Ladle Range,

Percent

4100 TYPE STEEL* 8600 TYPE STEEL*

AISI-41 18

Steel

Mn-Ni-Mo

Replacement

Mn-Mo

Replacement

AISI-8620

Steel

Mn-Ni-Mo

Replacement

Mn-Mo

Replacement

Carbon .18 - .23 .16 - .21 .16 - .21 .18 - .23 .16 - .21 .16 - .21

Manganese .70 - .90 1.00 - 1.30 1.00 - 1.30 .70 - .90 1.00 - 1.30 1.00 - 1.30

Chromium .40 - .60 r r .40 - .60 r r Nickel r .20 - .40 r o i o .40 - .70 r Molybdenum .08 - .15 .15 - .25 .25 - .35 .15 - .25 .25 - .35 .35 - .45

Dib**,min,in 1.15 1.25 1.25 1.40 1.55 1.50

Dlfc>**>rnax,in 1.75 1.85 1.80 2.35 2.30 2.20

Die**, min, in 3.15 3.15 3.30 4.30 4.40 4.45

* The carbon content of the Mn-Ni-Mo and Mn-Mo replacements for the 8620 and 4118 steels

is 0.02% lower than the standard grades. It is anticipated that this 0.02 % reduction in

carbon content would be maintained throughout the entire replacement series. ** Dj values calculated from chemistry +1/4 range.

r = residual level; Si range = .15-.35%; S = .05% max; P = .04% max.

CO

lil z

0

cr<

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o

DC

V

4

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1 1 8 H

A L C U BAND

A T E D H - B / N D -•••

1

\\

i

! V

rv

••X

!»>J A

V k.\

4 8 12 16 20 24 28 32

JOMINY DISTANCE , 1/16 INCH

Figure 3 - Standard Jominy hardenability band for AISI 4118

steel compared to the calculated hardenability band for the

chrome-free steels.

P30-14
EXPERIMENTAL RESULTS

Small, 100 lbs, experimental heats of the chromium-free replacement steels were

produced in order to experimentally confirm that hardenability, microstructure, and

properties predicted for the chrome-free steels correlate with measured values. The

experimental heats were produced at Climax Molybdenum and processed into forged

bars or plate samples. Jominy hardenability tests were performed on the base

composition, and carburized Jominy hardenability tests were performed to evaluate the

high carbon hardenability of the chromium-free steels. Fracture toughness tests on

specially prepared high carbon heats of the chromium-free replacement steels were

conducted in accordance with procedures outlined in ASTM E-399. Six experimental

heats were produced. The chemical analysis of the six heats is shown in Table VIII.

TABLE VIII - CHEMICAL ANALYSIS OF EXPERIMENTAL HEATS OF

REPLACEMENT STEEL FOR THE AISI 8600 SERIES

Chemical

Analysis

% by Weight

SAMPLE IDENTIFICATION

8620

Base Line

Heat

Mn-Ni-Mo

Lo-Side

Heat #41

Mn-Ni-Mo

Hi-Side

Heat #38

Mn-Ni-Mo

Hi-Carbon

Heat #39

Mn-Mo

MidHeat #47

Mn-Mo

Hi-Carbon

Heat #70

Carbon .21 .19 .23 .91 .21 .93

Manganese .85 .86 .97 1.21 1.12 1.14

Silicon .23 .18 .35 .32 .21 .31

Sulfur .03 .03 .03 .03 .03 .02

Phosphorus .02 .02 .03 .02 .02 .01

Nickel .47 .55 .68 .58 .04 .04

Chrome .46 .03 .11 .05 .05 .05

Molybdenum .20 .21 .26 .34 .30 .40

Aluminum .053 .071 .085 .09 .082

Nitrogen - .009 .015 .012 .019 .011

DIb, inch 1.75 1.35 3.05 - 1.75 -

Die, inch 5.15 3.95 7.25 6.35 5.10 6.20

Oominy hardenability test results conducted on the low side heat of the Mn-Ni-Mo steel

(heat #41) are shown in Figure 4. The measured curve falls below the standard curve and exhibits lower than the expected results. Investigation revealed that the grain size

in this heat (ASTM #9) was finer than anticipated. In general, in hardenability calculations an ASTM grain size of No. 7 is assumed. The measured curve is compared to

calculated Jominy curves assuming grain sizes of ASTM #7 and ASTM #9 in Figure 5. Note that the calculated hardenability curve assuming grain size of ASTM #9 falls

closer to the experimental curve.

P30-15
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ro

UJ§*

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oc

8

1 6 2 0 L 1

0-31

1 D E C A1 L C U L A T E D

H EAT# 4 1 ME a s u n —

A

1'

4 8 12 16 20 24 28 32

JO MIN Y DISTANCE . 1/16 INCH

Figure 4 - Hardenability curve measured for the Mn-Ni-Mo low

side heat #41 compared to the calculated 8620 low side

hardenability curve.

CO

CO

UJz

Q

cr<

io

UJ§

o

o

cr

c ALCU LATE

r~D , G S

-

-4

7

1 1 c ALCU

1 LATE

1 D . G 3

- 9 _ —

FIEAT# 4 1 M E A 3 U

-4-

\\

[%.

4 8 12 16 20 24 28 32

JOMINY DISTANCE , 1/16 INCH

Figure 5 - Comparison of the measured low side hardenability

to calculated hardness curves for different grain sizes.

P30-16
The 3ominy curve for the high side heat of the Mn-Ni-Mo steel is compared to the high

side curve for the standard 8620 steel in Figure 6. The chemical analysis of the high

side heat exceeds the desired analysis resulting in a higher hardenability and, hence, a higher measured curve. In producing experimental heats, it is very difficult to obtain

exactly the analysis desired.

3ominy hardenability tests conducted on the mid-band heat of the Mn-Mo chromiumfree steel are shown in Figures 7 and 8. This steel also exhibited a finer than expected

grain size, and Figure 8 demonstrates that when the calculation is adjusted to include

the finer grain size the measured and calculated curves correlate very well.

The occurrence of finer grain size in the experimental heats is related to a high

aluminum content (.07 - .09 percent) and a high nitrogen content, which provides a

larger number of AIN particles to affect the nucleation and growth of austenite.

CO

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-SIDE

8 M E A CAL

SURE

C U L A

D

TED —

\\\% N \

4 8 12 16 20 24 28 32

JOMINY DISTANCE , 1/16 INCH

Figure 6 - Comparison of the high side Mn-Ni-Mo heat #38 to

the calculated 8620 high side hardenability curve.

Carburized 3ominy hardenability tests were performed on the three sample heats and

results are shown in Figures 9 through 12. Figure 9 shows the 8620 standard steel base

line test. All tests were pack carburized in new carburizing compound and heated at

1700° F for eight hours, and direct quenched using a standard 3ominy bar quench

fixture. The surface carbon content in the chromium-free steels was in general

reduced, as compared to the 8620 standard sample. This reduction indicates a reduced

tendency to form carbides as a result of the elimination of chromium, and implies that

these steels would tolerate a greater variation in carbon potential in the furnace than

the standard chromium containing steels.

P30-17
05

cn

LUzQ

cc<

X

o

UJ§

O

o

cc

4 8 12 16 20 24 28 32

JOMINY DISTANCE , 1/16 INCH

Figure 7 - Comparison of the mid-band Mn-Mo heat #47 to the

calculated 8620 mid-band hardenability curve.

COCO

UJz

0

cc<1

o

UJ5

*o

o

cc

CALCULATED, GS-7

CALCULATED, GS-9-- -

! 1 l 1 l 1

,

HEAT# 4 7 M E A 3 U TED —

V\

-

4 8 12 16 20 24 28

JOMINY DISTANCE , 1/16 INCH

Figure 8 - Comparison of the measured hardenability were for

heat #47 to calculated hardness curves for different grain sizes.

P30-18
to

to

UJz

Q

£E<

I

O

UJ5

*o

o

tE

JOMINY DISTANCE , 1/16 INCH

Figure 9 - High carbon (carburized) hardenability curves developed in the 8620 base line steel.

JOMINY DISTANCE , 1/16 INCH

Figure 10 - High carbon (carburized) hardenability curves developed in the Mn-Ni-Mo high side heat #38.

P30-19
JOMINY DISTANCE , 1/16 INCH

Figure 11 - High carbon (carburized) hardenability curves developed in the Mn-Mo mid-band heat #47.

4 8 12 16 20 24 28 32

JOMINY DISTANCE , 1/16 INCH

Figure 12 - High carbon (carburized) hardenability curves developed in the low side Mn-Ni-Mo heat #41.

P30-20
Results of the fracture toughness studies performed on high carbon steels to simulate case microstructure are given in Table IX. High carbon chrome-free replacement steels

were poured especially for this program. The data for 8695 steel, included for

comparison, was obtained at IH under another in-house program. Kj[ c

values are averages of at least three measurements. X-ray retained austenite (RA) measurements

given in the table were made using a rotating and tilting staged) to compensate for the

texture in the specimens.

TABLE IX

PLANE STRAIN FRACTURE TOUGHNESS VALUES FOR HIGH CARBON STEELS

Sample %C Rc RA Kl c, Ksi /iru

Mn-Mo 0.93 60.5 22% 16.2

Mn-Ni-Mo 0.91 60.5 23% 16.1

8697 0.97 60.0 23% 15.2

Fracture toughness values for all three steels are essentially the same. A slightly lower

value for 8697 steel may be attributed to a slightly higher carbon content of 8697 steel.

COST BENEFIT OF THE CHROMIUM-FREE STEELS

The cost benefit to be derived from the development of chromium-free replacement

steels for the 8600 and 4100 standard steels may be estimated as the difference in the

costs required to purchase the chromium-containing standard steels vs. the costs

required to purchase the chromium-free replacement steels, assuming that a chromium

supply interruption has occurred in 1985.

§ 12.00 oW

10.00 —

« 8.00

O

o

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DCH

X

LU

111 Q

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cra

6.00 —

4.00 —

i r

8620 EXTRA /

/

• /

/

2.00 —'

'

“ /

/

^

4118 EXTRA _

19 6 4 19 7 2 19 8 0 19 8 8

TIME , YEARS

Figure 13 - Chemistrys grade extras for AISI 8620 and 4118 steels: 1960 to 1982

P30-21
§

$/c 2 4.00 1 1 1 r S C E N A R O—' -

d

2 0.00 _ H

(0 S C E N A R O c :

o

o

16.00

S C E N A R 1 0 3

' ^ c

EXTRA

12.00

8.00 SCENARIO A

'—

'

UJO

4.00 -

<

CL

1

1

1 1

i (D 19 8 1 19 8 3 19 8 5

TIME , YEARS

Figure 14 - Predicted effect of the various 1985 chrome supply

scenarios on the 8620 steel grade extra.

The cost of the two AISI standard grades of steels in terms of their grade extras is shown for the time period of 1960 to 1982 in Figure 13. The effect of the various

chromium supply scenarios on the 8620 steel grade is projected to 1985 in Figure 14.

The 8620 steel grade extra increases from $12.10/cwt in Scenario A to $22.40 in

Scenario D. The steel grade extras for the corresponding Mn-Ni-Mo and Mn-Mo

chromium-free replacement steels are shown in Figures 15 and 16 respectively. Both of

the replacement steels at a 1981 grade extra are more expensive than the AISI 8620

steel. In 1985 the Mn-Ni-Mo replacement steel becomes cost-effective only under the

1-5

o 2 4.00 —

i

1 1 r i

i—to

2 0.00

oo

16.00 SCENARIOS C&D ~ SCENARIO B * <

CL

h12.00

SCENARIO A X '

in 8.00 —

UJo

<

CL

4.00

1 1 1 1 1 0

1981 1983 1985

TIME , YEARS

Figure 15 - Predicted effect of the various 1985 chrome supply

scenarios on the Mn-Ni-Mo (replacement for 8620) grade extra.

P30-22
chromium supply Scenarios C and D. The Mn-Mo replacement for the AISI 8620 at a 1981 grade extra is slightly more expensive ($9.50/cwt vs. $9.60/cwt), but it becomes

cost-effective under all scenarios in 1985. Based on the 1978 domestic production of

2,042,176 tons of 8600 type steel, the Mn-Ni-Mo and Mn-Mo steels provide an annual

cost advantage of 257.3 and 322.6 million dollars, respectively, in Scenario D.

h

§

$/c 2 4.00 ( j

1 1 1

K

CO

2 0.00 ' —

o

O

16.00

SCENARIOS C & D .»•—' '—' " SCENARIO B j—

CLX '

12.00

SCENARIO A

UJ 8.00

LU 4.00 o

<

CE

1 1

1 1 1 1981 1983 1985

TIME , YEARS

Figure 16 - Predicted effect of the various 1985 chrome supply

scenarios on the Mn-Mo (replacement for 8620) grade extra.

A similar cost analysis for the AISI 4188 and its two chromium-free replacement steels

is shown in Figures 17 to 19. Projected 1985 steel grade extra cost for 4118 steel

increases from $5.95/cwt, Scenario A, to $15.00/cwt under the conditions of

Scenario D. These chrome-free steels, however, are not cost-effective in Scenario A,

H

§

O 2 4.00 r

1 1 r 1

2 0.00 - H

COO 16.00 _ — o S C E N A R 0 [) ^ ^' < 12.00 I GCH S C E N A R 1 0 C X

UJ 8.00 -

S C E N A R 1 0 3 i—— ^ - -

UJ 4.00 — —

D SCENARIO A

—* <

ccrn l 1 1 L 1 1981 1983 1985

TIME , YEARS

Figure 17 - Predicted effect of the various 1985 chrome supply

scenarios on the 4118 steel grade extra.

P30-23
B and C. Under the 1985 chromium supply Scenario D, the Mn-Ni-Mo and Mn-Mo

replacement steels for 4118 steel provide an annual cost advantage of $198.5 and

$233.8 million dollars, respectively. This cost savings is based on 1978 domestic

production of 2,687,512 tons of AISI 4100 series steel.

§

o

te

i-coo

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TIME , YEARS

Figure 18 - Predicted effect of the various 1985 chrome supply

scenarios on the Mn-Ni-Mo (replacement for 4118) grade extra.

5

TIME , YEARS

Figure 19 - Predicted effect of the various 1985 chrome supply

scenarios on the Mn-Mo (replacement for 4118) grade extra.

P30-24
The economic advantages of using chrome-free steels under various scenarios is summarized in Table X. The replacement of both 8600 and 4100 types steels with the

Mn-Ni-Mo and Mn-Mo steels, under conditions of Scenario D, will result in a combined

cost advantage of $455.8 and $556.4 million, respectively.

TABLE X - ECONOMIC ADVANTAGE OF CHROME-FREE STEELS

Grade Extra Cost for Each Scenario, $/cwt

Steel

Type

1981 1985A 1985B 1985C 1985D

AISI 8620 Steel 9.50 12.10 14.90 17.80 22.20

Mn-Ni-Mo 8600

Replacement 10.85 13.20 15.25 15.90 15.90

Mn-Mo 8600

Replacement 9.60 11.85 13.15 14.30 14.30

AISI 4118 Steel 4.90 5.95 7.65 10.50 15.00

Mn-Ni-Mo 4100

Replacement 7.50 9.20 10.25 11.30 11.30

Mn-Mo 4100

Replacement 7.10 8.80 9.80 10.65 10.65

Annualized Cost Advantage (Penalty)* - Millions of Dollars

Mn-Ni-Mo 8600

Replacement (55.2) (44.9) (14.3) 159.3 257.3

Mn-Mo 8600

Replacement (4.1) 10.2 71.4 224.6 322.6

Mn-Ni-Mo 4100

Replacement (139.8) (174.7) (139.7) (42.9) 198.5

Mn-Mo 4100

Replacement (118.3) (153.2) (115.6) (8.0) 233.8

Combined Cost Advantage (8600 and 4100 Steels) - Millions of Dollars

Mn-Ni-Mo

Replacement (195.0) (219.6) (154.0) 116.4 455.8

Mn-Mo

Replacement (122.4) (163.4) (44.2) 216.6 556.4

* Annualized cost advantages based upon 1978 domestic steel production of

2,042,176 tons of 8600 type steels and 2,687,51 2 tons of 4100 type steels.

P30-25
The Mn-Mo replacement steels provide approximately a 22 percent additional cost

advantage over the Mn-Ni-Mo steels. This is correct only under the assumption that the

molybdenum market will be capable of absorbing the additional increased demand

without increasing the molybdenum price above the critical ratio of 2.4. Based upon

the 1978 production volume of 4,727,000 tons (combined) of 8600 and 4100 series steels,

17.5 million lbs of molybdenum were required to produce the standard steels. In eliminating chromium to develop the new chromium-free replacement grades, an

increase in molybdenum content is required. The manganese-nickel-molybdenum steel

grade would require 27 million lbs of molybdenum to produce 4,727,000 tons of steel.

This is an increase of approximately 10 million lbs of new molybdenum demand. The

manganese-molybdenum steel grade would require 35 million lbs of molybdenum, an

increase of 18 million lbs of new demand. The selection of either the manganesenickel-molybdenum grade or the manganese-molybdenum grade as the most costeffective replacement for the 8600 and 4100 standard steels is dependent upon the

ability of the molybdenum market to accept and accommodate a rapid increase in new

domestic demand at the time the chromium supply interruption occurs. If the

molybdenum market is capable of accepting 18 million lbs of new demand without

raising the price of molybdenum above the critical ratio of 2.4, then the manganesemolybdenum replacement is the most cost-effective. Based upon current evaluation of

the molybdenum market, it seems that the molybdenum market will be able to absorb

the additional demand and the manganese-molybdenum replacement steel would be the

most cost-effective chromium-free replacement steel.

CONCLUDING REMARKS

Based upon the several socio-economic scenarios and their projected impacts on

alloying elements cost and availability in 1985, two chromium-free substitute steels

have been designed. Both steels, a manganese-molybdenum substitute and a manganese

nickel-molybdenum substitute, are expected to provide microstructure, heat treat

response, and mechanical properties equivalent to the 8600 and 4100 steels. The use of

chromium at the present time is highly cost-effective and the implementation of the

chromium-free steels under the present conditions would not be in the best economic

interest of the nation. However, the development of the two chromium-free steels

provides materials information which can be saved or stored in a "National Substitution

Data Bank" against a future time when an interruption in the chromium supply occurs. Under the worst case scenario - breakdown in South African chromium supply and at the

same time, cut-off of Soviet and Albanian sources - the use of chromium-free steels is expected to provide savings of over half a billion dollars annually. It is also possible

that the existence of chromium-free substitute steels could moderate future chromium

producer supply and pricing policies. Thus, the development of these substitute steels

significantly enhances the National "substitution preparedness" with regard to

chromium as used in the constructional alloy steels. Obviously, additional testing is required to develop engineering design data and to confirm equivalency of at least one

of these chromium-free steels under production conditions prevalent in U.S. industry.

Availability of this data will further encourage the ready adoption and widespread use by industry of the new chromium-free steel in the event of a shortage in the supply of

chromium.

P30-26
ACKNOWLEDGMENT

This project was sponsored by the United States Department of the Interior, Bureau of

Mines under Contract No. 30 113104. The support of Messrs. 3. T. Dunham and

K. W. Mlynarski of the Bureau of Mines, Washington Office, in developing the project is sincerely appreciated. The contract was administered through the Materials Section of

the Bureau of Mines, Albany Research Center, under the direction of Mr. H. W.

Leavenworth with Mr. Max L. Glenn serving as technical project officer. The

assistance of Mr. Glenn in providing technical and administrative assistance is gratefully acknowledged.

REFERENCES

1. Morning, 3. L., Mineral Commodity Profiles, Chromium - 1977, Bureau of Mines,

U.S. Department of Interior, May 1977.

2. Contingency Plans for Chromium Utilization, NMAB Commision on Sociotechnical

Systems, National Research Council, National Academy of Sciences, Washington,

D.C., 1978 - NMAB - 335.

3. Breen, D. H., Walter, G. H., Keith, C. 3., and Sponzilli, 3. T., "Computer Based

System Selects Optimum Cost Steels," Metal Progress, November 1973.

4. Keith, C. 3., Sponzilli, 3. T., Sharma, V. K., and Walter, G. H., "International

Harvester's CHAT System for Selecting Optimum Compositions for Heat Treated

Steels," Hardenability Concepts with Applications to Steel, edited by Doane, D. V.

and Kirkaldy, 3. S., published by AIME, 1978.

5. Ralph, W. L., "Lessons from Nickel," Molybdenum Supplement, American Metal

Market, 3uly 1982.

6. U.S. Bureau of Mines, Mineral Commodity Summaries 1980, Washington, D.C.,

Government Printing Office, 1980.

7. Shin, S. W., and Sharma, V. K., "Application of A Tilting and Rotating Specimen

Stage to X-Ray Retained Austenite Measurement in Textured and Coarse Grained

Steels," SAE Trans., Paper 800428.

P30-27
.

.

.
CHROMIUM-FREE STEELS FOR CARBURIZING

D. E. Diesburg, G. T. Eldis and H. N. Lander*

INTRODUCTION

Despite the fact that most commonly used carburizing steels contain on

the order of 0.5 to 1% chromium, the prospect of a restriction or total cut-off

of chromium supply is not terribly troubling to the carburizing steel metallurgist. There already exist several standard chromium-free carburizing grades.

The general methodology of developing new carburizing steels to replace existing

grades, substituting more abundant and less expensive alloying elements to

reduce overall steel cost while maintaining desired properties, is well

developed, thanks to past materials shortages of one kind or another. And

especially over the last five years or so, we have gained a great deal of

knowledge concerning the effects of all the commonly used elements, individually and in combination, on the critical properties of carburized components, so that designing chromium-free steels to meet the requirements of

typical carburizing steel applications should pose very little difficulty.

Indeed, the only area in the realm of carburizing steels where the prospect

of a chromium shortage gives us pause is the relatively specialized application

of elevated temperature service. At present, all elevated temperature carburizing steels contain chromium, and the steel properties are heavily dependent

on a dispersion of alloy carbide precipitates. However, the authors are confident that, if work is begun now, we can be ready to meet any future chromium

shortage with a suitable low-chromium or chromium-free steel for this application, too.

*D. E. Diesburg, G. T. Eldis and H. N. Lander are Research Supervisor, Research

Manager, and Sr. Vice President, Research and Development, respectively. Climax

Molybdenum Company of Michigan, Division of AMAX of Michigan, Inc., Ann Arbor,

Michigan.

P31-1
In the discussion which follows, we will first look, at existing standard

carburizing steel grades to get an indication of the economic impact of a

chromium shortage. We will then consider the effects of various alloying

elements on hardenability and on the mechanical properties critical to carburizing

steel performance, to show the directions we are likely to follow in our pursuit

of new chromium-free carburizing steels and the properties these steels are

likely to possess. Lastly, we will look briefly at the most critical problem

for chromium-free carburizing steels, namely, elevated temperature service.

EXISTING STANDARD GRADES

In any discussion of carburizing steels, we usually find it convenient to

classify the steels into one of three core hardenability ranges, low, medium or

high. Figure 1 illustrates these three ranges as schematic hardenability bands

showing the depth of hardening in an end-quench test to be expected for each

range. Table I catalogues various domestic (SAE standard) and European grades

in the respective groupings.

The low-hardenability grades consist of those steels used for relatively small carburized components, specified by the U.S. automotive industry

as SAE 4000, 4100, 4400, 4600, 5100, 8600 and 8700. The SAE 8600 and 8700

grades have the highest alloy hardenability and represent the hardenability most

often required for such components. SAE 4028 has a hardenability similar to

that of SAE 8617; however, a substantial amount of the hardenability comes from

the carbon content. This is an important point. Carbon is a very potent element

in increasing hardenability, and as such is about the most cost-effective

P31-2
"alloying" element. But higher core carbon contents generally lead

1 2 to relatively inferior bend strength properties. ’ Thus, the SAE 4028

grade would be used primarily in those applications where specific bending

stresses are relatively low. Of the low-hardenability grades, SAE 4000, 4400

and 4600 are chromium-free steels. The other grades contain nominal chromium

contents ranging from 0.5 to about 1.0%. In general, the alloy cost of the

chromium-free grades is higher.

The medium-hardenability steels are represented by grades such as SAE 4300

and 8800, both containing 0.5% Cr. These grades are used when the section sizes

being quenched are too large for the low-hardenability steels to maintain an

adequate core hardness. At present, there are no standard chromium-free steels

that fall into the medium-hardenability classification.

The high-hardenability grades include such steels as SAE 4815, 3310 and

9310. The SAE 4800 grade is a chromium-free Ni-Mo steel while the SAE 9310

and 3310 steels are Ni-Cr steels containing 1.2 and 1.6% Cr, respectively.

These steels are used to produce large components with tough martensitic

cores and also for smaller components expected to encounter service

too severe for the low- or medium-hardenability steels. In this respect, these

grades are often considered to be high performance grades.

Within the various core hardenability ranges, there are several new grades

which have been proposed and made commercially available which represent more

economical alloy compositions than existing standard grades. These proposed

alternates have been assigned SAE EX (or PS) numbers. Table II shows a few

examples of EX steels designed by the Climax Molybdenum Co. and compares them

with the SAE grades they were designed to replace. All of these alternate

P31-3
steels contain higher levels of chromium, manganese and molybdenum than the

corresponding SAE grades, illustrating a principal reason for the frequent use

of chromium in carburizing steels: At today's availability and price, chromium

provides an economical source of alloy hardenability

.

But what of the technical merits of chromium? Cost factors aside, can

we achieve desired properties without this alloying element? Let us consider

various attributes of carburizing steels and what we know about the effects

of alloy additions.

HARDENABILITY

We normally subdivide hardenability into two categories. Core hardenability

is a measure of the ability to achieve hardness at depth in the low carbon core

of the carburized component. It is the hardenability of which we've been

speaking in the discussion thus far of hardenability ranges. Case hardenability

is a measure of the ability to achieve martensitic microstructures, free of

softer pearlitic and bainitic transformation products, in the high carbon case

of the carburized component.

Figure 2 illustrates the effects of the common alloying elements on core

3 hardenability. Molybdenum and manganese are, individually, of comparable effectiveness, although molybdenum is substantially more effective if the steel

*There is no standard grade with a hardenability range similar to EX55.

Its hardenability exceeds that of SAE 9310 and SAE 3310, but with a nominal

carbon content of 0.17% and a rather low chromium level of 0.55%.

P31-4
contains a minimum of about 0.75% Ni. Chromium is somewhat less effective

than manganese or molybdenum, and nickel is relatively ineffective by itself.

Silicon has no effect on core hardenability at the levels typically found in

carburizing steels. If we included a cost factor into these data, turning the

abscissa into "alloy cost," the order of the elements would be reversed with

manganese being by far the most effective alloying addition, followed by chromium, molybdenum and nickel. That is at current prices and availability, of

course. From the core hardenability standpoint, elimination of chromium poses

no major difficulties. The hardenability decrease resulting from elimination

of, say, 1% Cr from the steel could be readily recovered by an increase of

0.6% in the combined Mn + Mo content. The ratio of manganese to molybdenum

selected for such a substitution would depend upon several factors, among them

the overall level of those elements already present in the steel, since excessive use of any one element can aggravate the problems of alloy segregation

during solidification. Other factors to be considered will be discussed later.

Figure 3 shows the influence of molybdenum, manganese, chromium and nickel

on case hardenability, in the carburized areas of the component where carbon

4

content is usually on the order of 0.8 to 0.9%. The relative effectiveness

of the elements is quite different than for core hardenability. In addition,

molybdenum and nickel share a common attribute which makes them, relative to

manganese and chromium, even more effective than Figure 3 implies. These

elements are immune to depletion from the matrix by the mechanism of surface

oxidation. During normal carburizing in endothermic atmospheres, the oxygen

potential is high enough to induce oxidation of chromium, manganese and

silicon.^ These elements are depleted from the matrix at the surface of the

P31-5
specimen, as illustrated in Figure 4 which shows the results of electron probe

6

microanalysis on a 1 . l%Mn - 1 . l%Cr steel carburized in an endothermic atmosphere.

Figure 5 shows the microstructural effects of this surface oxidation, a decrease in hardenability at the surface and transformation to pearlite and

bainite rather than martensite. Depending on depth of oxidation/depletion,

size and quenching rate of the component, subsequent surface treatment (lapping,

shot peening) and in-service loading experienced by the component, these nonmartensitic transformation products at the surface could ultimately prove detrimental to performance.

Turning back to Figure 3 and our example of removing 1% Cr from the steel

composition, we could maintain the same level of case hardenability by simply

adding 1% Mn. (Note that this is substantially more than the 0.6% Mn required

to maintain a constant level of core hardenability, as discussed before.)

However, given that the steel will most probably already contain at least

0.5% Mn, a full manganese-for-chromium substitution would not be a good choice.

It would result in an "excessive" level of manganese (1.5% or more), and would

only aggravate the surface oxidation problem (Figure 4) . A more suitable substitution might be, say, 0.4%Mn + 0.2%Mo, which would maintain both the case

and core hardenability lost by the removal of 1% Cr. Regardless of the exact

combination of elements chosen to substitute for the chromium, it is quite

clear that the substitution can be made without experiencing any reduction of

steel performance from the hardenability standpoint.

P31-6
MECHANICAL PROPERTIES

There are two rather independent mechanical properties of carburizing

steels that determine their overall performance: (1) the ability to resist

pitting and (2) the ability to resist bending. In most applications, carburized components are overdesigned with respect to bending fatigue simply

because complete fracture is a less tolerable situation than is surface

degradation such as pitting, although severe pitting also can eventually lead

to complete fracture.

Pitting Fatigue

Factors influencing pitting fatigue are more difficult to evaluate

experimentally than are the factors influencing bending fatigue. Although

it is generally believed that pitting behavior is dependent on hardness and

microstructure, at times it has been difficult to confirm this experimentally . ^

9 Recent research has shown pitting to be associated with inclusions. Pitting

resistance has also been shown to be dependent on the lubricant.^ A proper

lubricant can maintain a film between the contact surfaces and significantly

improve the pitting fatigue lives of carburized components. Soft microstructures are generally accepted as having poor resistance to pitting, although

hard microstructures can also exhibit pitting. Surface decarburization, the

presence of a large amount of retained austenite, and the transformation of

austenite to pearlite or bainite during quenching can cause soft regions to

occur in a carburized case and thus are believed to lead to an eventual pitting

failure. Figure 6 shows pitting damage as revealed in the scanning electron

microscope

.

P31-7
All common alloying elements used in carburizing steels, namely chromium,

manganese, nickel and molybdenum, are added to obtain the required hardenability

In this respect, the elements are interchangeable as already discussed. However

there are independent limiting factors that determine the maximum amount of

each element that should be added, in addition to the segregation and surface

oxidation problems already mentioned, which could affect pitting resistance.

Nickel and manganese are both austenite stabilizers and therefore must be

used in moderation if retained austenite in the carburized case is expected to

be a problem. Chromium also has a carbide formation limitation. High chromium

contents are believed to promote the formation of large carbides in the case

during carburizing, thus steels such as SAE 3310 and SAE 9310 generally require

extra care during carburizing and heat treatment.

Reducing or eliminating the chromium contents in carburizing steels is

not expected to have a strong effect on the pitting resistance, provided the

corresponding loss in hardenability is regained by adding manganese, nickel

or molybdenum and retained austenite content remains under control. In support

of the contention that chromium-free steels will offer adequate pitting resistance compared to chromium-containing steels, one can look at the current

successful use of existing grades of chromium-free steels such as SAE 4000,

4400, 4600 and 4800. There are no reported indications that these grades

are any more prone to pitting than the chromium-containing steels.

Bending Fatigue and Impact

Bending fatigue properties of carburized steels are highly dependent on

the compressive residual stress that develops at the carburized surface during

P31-8
11 12 quenching. ’ This is illustrated nicely by the improvement in high cycle

13 fatigue limit resulting from shot-peening, shown in Figure 7. Shot-peening

prior to carburizing does not provide a similar improvement in fatigue behavior

14 (Figure 8). Altering the residual stress at the surface so as to make it

less compressive, as by improper heat treatment or other processing, can have

a dramatic negative effect on fatigue behavior. This is illustrated in Fig13 ure 9 by the data for EX55 subjected to a carbonitriding treatment. The

particular cycle employed resulted in a very high (70%) retained austenite

content at the specimen surface. Despite this high austenite content, the

bulk residual stress at the surface was compressive and the fatigue properties quite good. A refrigeration treatment reduced the austenite content to

40%, but also resulted in a net bulk tensile residual stress at the surface.

The effect on fatigue properties was disastrous.

The only expected effect of chromium per se on constant load amplitude

high-cycle fatigue behavior would be the result of chromium depletion from

the surface, either by severe oxidation or by massive carbide formation. This

could allow pearlite or bainite to form at the surface, thus causing the surface

to have less compressive residual stress, perhaps even a tensile stress.^

Figure 10 illustrates the theoretically calculated effect of compressive residual stress on fatigue limit for various assumed degrees of surface oxida16

tion. Here, the oxidation product is assumed to act as an internal defect

(stress concentrator) which intrinsically reduces fatigue limit. The importance of maintaining a substantial compressive residual stress at the surface is quite clear. Reducing or eliminating chromium in carburizing steels

is thus expected to have no detrimental effect on high-cycle fatigue behavior.

P31-9
provided the lost case hardenability is replaced with other alloying elements

to ensure the surface microstructure remains martensitic. Indeed, to the extent

that chromium-free alternate steels might be less susceptible to surface oxidation effects, an improvement in properties might be expected.

The apparent sole dependence of bending fatigue performance (fatigue

limit) on the surface residual stress is true when the applied loads are

rather low, with maximum loading in the vicinity of the laboratory-measured

fatigue limit. However, most carburized parts in service are subjected to a

wide spectrum of loading, and the infrequent occurrence of loads well in excess

of the fatigue limit must be anticipated. Such overloading can result in

material damage, so that the effective fatigue limit after the overload is

reduced .The magnitude of overload that can be tolerated before significant damage occurs can simply be referred to as "toughness," and is^ highly

dependent on alloy content as well as the fracture toughness and residual

14

stress further away from the carburized surface. The relative ability of

the steels to resist overloading can be evaluated by measuring the impact

fracture strength of the carburized component, that is, the load required to

fracture the component in a single blow. Climax has been evaluating the effect

of various alloy combinations on the impact fracture strength for several years

and has published some generalizations that can be used to assess the contribution of chromium to bending fatigue life when random high and low loadings are

. 18 encountered

.

For low-hardenability steels. Figure 11 indicates that chromium generally

has a negative effect on impact fracture strength over the range of 0.2 to 1.2%.

Thus, no detrimental effect on impact fracture strength or "toughness" is

P31-10
expected if chromium is eliminated from carburizing steels, provided adequate

alloy substitution is made to maintain hardenability

.

In the medium-hardenability range, the steel compositions generally include various nickel-chromium-molybdenum combinations. The influence of chromium

on impact fracture strength is dependent on the nickel content of the steel.

Combined with nickel contents of 1% or less, chromium shows a negative effect

on impact fracture strength similar to that experienced in the low-hardenability

steels. In combination with nickel contents exceeding 1.0%, chromium has no

influence. So again, using chromium-free steels for medium-hardenability should

have no detrimental effect on "toughness."

All high-hardenability steels contain more than 1.5% Ni, and as might be

expected from the trends in low- and medium-hardenability steels of various

chromium and nickel contents, chromium content has no effect on impact fracture

strength in the high hardenability steels. The excellent performance of

SAE 4800 steels provides ample evidence of the toughness of chromium-free

high-hardenability steels. EX55 contains 0.55% Cr, and it has exhibited even

better performance than SAE 4817 (other factors, such as the higher hardenability

of EX55, are partly responsible for this). A chromium-free version of the EX55

composition, containing 2% Ni and 1% Mo, has been evaluated in the impact fracture test and has exhibited similar behavior to that of EX55. Therefore, in

all three hardenability ranges, chromium-free steels are expected to have a

toughness equal to or better than currently used chromium-containing steels.

P31-11
ELEVATED TEMPERATURE SERVICE

Special steel compositions are available for carburized components to be

used at temperatures up to 315 C (600 F) . Table III gives some examples. Such

steels are required to retain adequate case hardness and toughness at these elevated service temperatures . Typical requirements might include a minimum case

hardness of 58 HRC after 1000 hours exposure at 315 C (600 F) , along with a

minimum fracture toughness specification.

To a large extent, these properties are dependent upon precipitation of

various alloy carbides during tempering at temperatures above the intended

service temperature, to form a strengthening dispersion stable at the operating

temperature with a morphology not overly detrimental to fracture toughness.

A variety of alloy carbides have been identified after heat treatment of

such materials, including M5 C in the lower carbon core regions of the

samples, and M23C 5 , M2 C and MC in the carburized case, depending on the exact

19 steel composition and heat treatment temperatures used. The extent to which

chromium enters into these various carbides is not well defined, and its necessity from the standpoint of precipitation kinetics and precipitate stability

is even less well defined. Thus, at present, we can't be certain how successful we would be in meeting the requirements of elevated temperature service

with a chromium-free composition. The authors are optimistic, however. From

2C

the standpoint of "carburizability , " the lower the chromium content the better.

Carbide precipitation and secondary hardening certainly do occur in vanadium,

21

molybdenum and tungsten alloyed steels without chromium additions, and such

precipitation is the major prerequisite for good elevated temperature properties. The comparably good performance of X-53 and X-2M, with 1 and 5% chromium.

P31-12
respectively, also bodes well for the development, if eventually necessary,

of chromium-free compositions. CBS 1000, with only 1% Cr, has performed even

better in Climax laboratory tests than X-53 and X-2M, and X-ray diffraction

studies have shown it contains no M23 C 6 , the chromium-rich carbide. Indeed,

the main difficulty in developing chromium-free steels for elevated temperature

service may have nothing to do with service performance at all, but rather

be related to slight atmospheric corrosion of components in inventory. Rather

remarkably, it was found in developing the CBS 1000 steel that nominally 1% Cr

was the minimum amount necessary to avoid problems of component degradation

19 in storage.

SUMMARY

Although chromium is a common alloy addition to carburizing steels, all

the evidence indicates it is not necessary for performance in typical (ambient

temperature) carburizing steel applications. Chromium is currently added to

low-, medium- and high-hardenability steels for economic reasons. Chromiumfree steels are expected to perform equally well and, in many situations, may

even provide better performance than currently used <_h romium-alloyed grades.

Additional work is necessary to determine the lull potential of chromium-free

carburizing steels to meet the requirements oi elevated temperature service,

but the prospects appear promising.

P31-13
REFERENCES

1. V. S. Sagaradze, "Effect of Carbon Content on the Strength of Carburized

Steel," Metal Science and Heat Treatment, March 1970, No. 3, pp. 198-200.

2. C. Kim and D. E. Diesburg, "Fracture of Case-Hardened Steel in Bending,"

submitted for publication in the Journal of Engineering Fracture Mechanics.

3. C. A. Siebert, D. V. Doane and D. H. Breen, The Hardenability of Steels , American Society for Metals, Metals Park, Ohio (1977), p. 101.

4. C. F. Jatczak, "Hardenability in High-Carbon Steels," Met. Trans., (1973),

pp. 2267-2277.

5. I. Kirman, G. Mayer and F. W. Strassburg, "Einfluss von Nickel and Randge

fiige auf die Brucheigenschaf ten einsatzgeharteter Stahle," Harterei-Techn.

Mitt. 29 (1974), pp. 88-94.

6. Y. E. Smith and G. T. Eldis, "New Developments in Carburizing Steels,"

Met. Engr. Quarterly, 16_, No. 2 (1976), pp. 13-20.

7. J. S. Learman and G. T. Eldis, "Effect of Bainite in the Outer Carburized

Case on Rolling Contact Fatigue Life," SAE Paper No. 760665.

8. S. L. Rice, "Pitting Resistance of Some High Temperature Carburized Cases,"

SAE Paper No. 780773.

9. D. H. Breen, formerly with International Harvester, currently with the

Gear Research Institute, unpublished research.

10. J. P. Sheehan and M. A. H. Howes, "The Role of Surface Finish in Pitting

Fatigue of Carburized Steel," SAE Paper No. 730580.

11. D. A. Sveshnikov, I. V. Kudryartsev, N. A. Gulyaeva and L. D. Golubovskaya,

Chemicothermal Treatment of Gears," Metal Science and Heat Treatment,

No. 7, July 1966, pp. 527-532.

12. Chongmin Kim, D. E. Diesburg and G. T. Eldis, "Effect of Residual Stress

on Fatigue Fracture of Case-Hardened Steels," ASTM STP 776 (1981)

pp. 224-234.

13. Chongmin Kim, D. E. Diesburg and R. M. Buck, "Influence of Sub-Zero and

Shot-Peening Treatments on Impact and Fatigue Fracture Properties of

Case-Hardened Steels," Journal of Heat Treating (3), 1 (1980), pp. 3-13.

P31-14
REFERENCES (Continued)

14. R. W. Buenneke, C. R. Dunham, M. P. Semenek, M. M. Shea, M. B. Slane

and J. E. Tripp, "Gear Single Tooth Bending Fatigue Test," SAE Paper

No. 821042.

15. B. Hildenwall and T. Ericsson, "Residual Stresses in the Soft Pearlitic

Layer of Carburized Steel," Journal of Heat Treating (3), 1 (1980),

pp. 3-13.

16. D. E. Diesburg, C. Kim and W. Bulla, "Impact and Fatigue Fracture of

Carburized Cases Related to Fracture Toughness and Residual Stress,"

Presented at Heat Treaters Conference, Wiesbaden, West Germany, 1981,

proceedings to be published.

17. T. B. Cameron and D. E. Diesburg, "The Significance of Impact Fracture

Strength of a Carburized Steel," to be presented at the Spring Meeting

of TMS-AIME in Atlanta, March 6, 1983 (to be published).

18. D. E. Diesburg and Y. E. Smith, "Fracture Resistance in Carburizing

Steels Part II: Impact Fracture," Metal Progress, 115 (6), June 1979,

pp. 35-39.

19. C. F. Jatczak, The Timken Co., private communication.

20. C. F. Jatczak, "Specialty Carburizing Steels for Elevated Temperature

Service," Metal Progress 113 , No. 4 (1978), pp. 70-78.

21.

G. A. Roberts, J. C. Hamaker and A. R. Johnson, Tool Steels , American

Society for Metals, Metals Park, Ohio (1962), pp. 215-218.

P31-1S
Table I

COMPOSITION, HARDENABILITY AND COST OF

COMMERCIALLY AVAILABLE STANDARD

CARBURIZING STEELS

Nominal Composition, Wt-%

Hardenability,b

Steel* C Mn Cr Ni Mo D|, mm S/ton'

Low Hardenability

4028 0.28 0.80 — — 0.25 40 157

4118 0.20 0.80 0.50 — 0.10 36 98

4422 0.22 0.80 — — 0.40 39 184

4617 0.17 0.55 — 1.80 0.25 37 323

5120 0.20 0.80 0.80 — — 41 55

8620 0.20 0.80 0.50 0.55 0.20 48 189

8720 0.20 0.80 0.50 0.55 0.25 50 208

*18CrMo4 0.18 0.65 1.05 — 0.25 53 163

*20NiCrMo2 0.20 0.75 0.50 0.55 0.20 45 189

*20MoCr4 0.20 0.75 0.40 — 0.45 48 253

*16MnCr5 0.16 1.15 0.95 - — 52 97

Medium Hardenability

4320 0.20 0.55 0.50 1.80 0.25 60 327

8822 0.22 0.85 0.50 0.55 0.35 62 252

High Hardenability

4815 0.15 0.50 — 3.50 0.25 66 512

3310 0.10 0.55 1.60 3.50 — >127 499°

9310 0.10 0.55 1.20 3.25 0.10 89 530“

*14NiCrMo13 0.14 0.45 0.95 3.25 0.25 99 559

*17CrNiMo6 0.17 0.50 1.65 1.50 0.30 >127 431

*20NiMoCr6 0.20 0.75 0.40 1.60 0.45 87 436

aAsterisk denotes European standard. Others are SAE grades.

Calculated from composition assuming ASTM grain size No. 7

cGrade extra charge based on U S. industry pricing.

Coes not include $25/ton electric furnace extra, which is required.

P31-16
Table II

COMPOSITION, HARDENABILITY AND COST OF SEVERAL

SAE GRADES AND THEIR EX ALTERNATES

Nominal Composition, Wt-%

Hardenability,"

Steel C Mn Ni Cr Mo D|, mm $/ton'

SAE 8620 0.20 0.80 0.55 0.50 0.20 48 189

EX 24 0.20 0.87 — 0.55 0.25 48 154

SAE 4320 0.20 0.55 1.85 0.50 0.25 60 327

EX 29 0.20 0.87 0.55 0.55 0.35 63 256

SAE 4815 0.15 0.50 3.50 0.25 66 512

EX 30 0.15 0.80 0.85 0.55 0.52 80 357

EX 55 0.17 0.87 1.80 0.55 0.75 >127 528

Calculated from composition assuming ASTM grain size No. 7.

bGrade extra charge based on U.S. industry pricing.

P31-17
Table III

CARBURIZING STEELS FOR

ELEVATED TEMPERATURE SERVICE

Nominal Composition, Wt-%

Steel C Mn Cr Ni Mo W V Cu

CBS 1000 0.14 0.5 1.15 2.9 4.7 — 0.3 —

X-53 0.18 0.4 1.05 2.1 3.3 - — 2.0

X-2M 0.15 0.3 4.90 — 1.4 1.4 0.4 —

P31-18
OISTANCE

FROM

QUENCHED

END

OF

SPECIMEN

IN

MILLIMETERS

31V0S SS3NQUVH 0 H3M>100H

u

in

cuo

00

•H

P31-19

DISTANCE

FROM

QUENCHED

END

OF

SPECIMEN

IN

SIXTEENTHS

OF

INCH
Figure 2 Effect of Various Elements on Hardenabi li ty at

0.2% Carbon Content (Reference 3)

P31-20
MULTIPLYING

FACTOR

i

i

Figure 3 Effect of Various Elements on Hardenabil ity at

0.85% Carbon Content (Reference 4)

P31-21
CASE DEPTH, mm

Figure 4 Chromium and Manganese Concentration Profiles Developed in

the Specimen Case During Carburizing of a 1.1 %Mn - l.l%Cr Steel (Reference 6)

P31-22
P31-23
4

TOOTH

ROOT

LINE OF INITIAL

MATING CONTACT

PITTING ZONE

Figure 6 Scanning Electron Micrographs of the Heavily Pitted Region

on the Compression Face of a Driven Gear

P31-24
STRESS,

MPa

2000

1000

0

Figure 7

DISTANCE FROM SURFACE , in.

DISTANCE FROM SURFACE, mm

0

-20

-40

-60

-80

M

CO

<0

Ul

a:h

(/) <

O

<0

UJ

a

300

«

200

*

co

(/>

UJ

a:

100 (A

0

Residual Stress Profiles and Stress-Life Curves for Carburized

SAE 4028, Before and After Shot-Peening (Reference 13)

P31-25
qi £oi

4 a von

iO

if)

LU

O

>-

o

P31-26

Figure

8

Load-Life

Data

for

Carburized

8620,

With

and

Without

a

Shot-Peening

Treatment

Prior

to

Carburizing

(Reference

14)
o

CL

2

CO

CO

UJ

och

CO

o

co

UJ

o:

DISTANCE FROM SURFACE t in.

0 O.Ol 0.02 0.03 0.04 0.05

200 F 1 1 ^ 1 1

CARBONITRIDED SAE EX55

TREATED AT

-85C (-120F)

-200

-400

NO SUBZERO

TREATMENT

1 ± 1 ±

20

-20

-40

-60

0 0.2 0.4 0.6 0.8 1.0 1.2

DISTANCE FROM SURFACE , mm

co

CO

UJ

CKh

CO

-I

<

3

o

co

UJoc

Figure 9 Effect of Refrigeration Treatment on Residual Stress and Stress-Life

Behavior of Carbonitrided EX55 (Reference 13)

P31-27
Figure 10 Predicted Effect of Surface Oxidation on Bending Fatigue Limit

for Various Levels of Residual Strength (Reference 16)

P31-28
IMPACT

FRACTURE

STRESS,

MPa

Figure 11 Influence of Chromium and Nickel Contents on Impact Fracture

Strength of Carburized Steels (Reference 18)

P31-29

IMPACT

FRACTURE

STRESS,

ksi

Report from Nachang Gear Plant, Beijing, China

RARE EARTH BORON STEEL (25 Mn-Ti-B) GEARS

Presented byDale H. Breen and Allen G. Gray

Rare Earth Boron Stee (25 Mn-Ti-B) is a chromium free carburized

steel developed jointly by 20+ plants, research institutes and

colleges in China. This new steel is based on utilizing the

material that is abundant in China. The outstanding features of

this steel are that it has good properties, is based on mature

technology, and has high reliability and low cost.

I. The Chemical Composition, Fundamental Properties and the

Refining of Rare Earth Boron Steel.

1. The chemical composition of rare earth boron steel in

weight percentage.

c Mn Si Ti B

Rare

Earth

0.22 1.30 0 . 20 0.06 0.001

- 20.05 0.28 - 1.60 - 0.45 - 0.12 - 0.004

2. The mechanical properties of the rare earth boron steel.

Heat

Treat

Tensile

Strength

(kg/mm z

)

Ultimate

Yield ? (kg/mm )

Elongation

m

Reduction

of Area

m

Impact

Strengt^

(kg .m/cin ) 850 C

Oil

Quench

200 C

Anneal

2 140 z? 110 2 10 2 40 2

6

3. The refining of the rare earth boron steel. In the past

10 years, steel plants throughout China such as Dayie

Steel Plant, Nanchang Steel Plant, and Peking Specialty

Steel Plant, have produced over 10,000 tons of rare earth

boron steel. Among those plants, the Dayie Steel Plant

has the highest capacity for producing the rare earth

boron steel.

P32-1
Rare earth boron ntccl can be refined by using electric

furnace, a combination of electric furnace and converter,

or simply converter. Regardless of the refining method,

the processing is simple, the quality of the product is

good, and yield is high (according to the tabulation by

Dayie Steel Plant, the yield by using any of the three

refining methods is about 97%). Therefore, this type of

steel can be produced in many steel plants in large

quantity.

4. Cost. According to the calculation made by Davie Steel

Plant, the price of the different kinds of rare earth

boron. steels when compared to 18 chrome-manganesetitanium steel is about 140 yuan "lower per ton. That

is to say the price of rare earth boron steel has a cost

of about 15% lower than the 18 chrome-manganese-titanium

steel. Therefore, wide application of rare earth boron

steel can lead to lowering the cost of production of

finished product.

II. Review of Plant Experiment and .Application of Rare Earth

Boron Steel.

Our plant has participated in the development, trial application

and promotion of the rare earth boron steel since 1564. Starting

November of 1974, wc have been substituting 18 chromc-manganesctitanium steel and 22 chrome-mangencse-moly steel with the rare

earth boron steel in the manufacturing of shift gears for

tractors and automobiles. In October of 1975 in a meeting called

"The 25 Manganese Titanium Boron Rare Earth Steel Application,

Exchange, Promotion and Testing" sponsored by the Ministry of

Metallurgical Industries and the Number One Ministry of

Machinery, we have reported in conjunction with other

P32-2
organizations the complete technical evaluation of the rare earth

boron steel. After that, we have expanded the application of the

rare earth boron steel to other types of complicated gears for

tractors and automobiles. Thus the rare earth boron steel has

•now become one of the major carburized gear steel in our plant.

To date, we have applied the rare earth boron steel in the making

of 76 types of gears in 7 different types of machineries.

Machine

Designation

45(55) 54 27 12 HJ-130 CA-10B 750

Number

of gears 32 2 22 6 4 6 4

From 1966 to September 1981, we have used over 12,000 tons of

rare earth boron steel. If we take the price difference between

rare earth boron steel and the chrome steel of 140 yuan, we have

saved so far in the cost of material alone about 1.8 million

yuan.

III. The Technical Properties of the Rare Earth Boron Steel.

1. The production of forged gear blank.

a. In the production of this material no matter whether

saw or shear is applied, there is no abnormal

phenomena. The properties are comparable with 18

chrorac-manganese-titanium and 22 chrome-manganesemoly steel.

b. The forging temperature range of rare earth boron

stool is 1250*0 maximum and 900**C minimum. Because

of the good hot workability of the steel, it is easy

to form in a closed die forging operation. The

oxide layer on the forged blank is thicker but it

is also fragile and loose. Therefore it is easy to

remove

.

P32-3
c. Normalization temperature for the forged gear blank

is 960°C. Usually we use air cooling. After

normalization, the hardness is IIB156 to IIB207 and 80%

of the blank has a hardness of about HB163 to 187.

2. Machineability. The machineability of the rare earth

boron steel. is good. The chips are easy to break and

there is no tangling phenomena. The machineability is

comparable with chrome steel. The cleanness of the

product after the machining operation is similar to those

of the chrome steel and can therefore satisfy the product

design demand completely. In our production, it was

discovered that when machining the center hole of a gear

blank of rare earth boron steel the machining speed has.

to be reduced. This could be due to the higher strength

and higher toughness of this particular steel.

3. Heat Treating of Gears. The rare earth boron steel has

a structure of fine grain and a direct quenching process

can be used after carburization. The features of the

process are as follows.

a. Carburization Speed. The speed of carburization fer

both rare earth boron steel and chrome steel is

basically the same. There is no difference in the

carburization production rate.

b. The Carbon Concentration in the Carburized Layer.

The dif fusivity of carbon in rare earth boron si tool

is lower than 18 chrome-mangancsc-titanium steel,

therefore there is a more uniform concentration of

carbon in the carbon diffusion layer compared to

chrome stool. Therefore better mctailographic

structure could be obtained.

P32-4
c. Grain Size Classification. After carburization and

quenching, 82% of the Martensite structure on the

surface diffused layer has a grain size of 2-4, the

other being 1 and 5. 86.5% of the residue Austenite

structure has a grain size of 2-4 and the remaining

being 1 and 5. 98% of the core Ferritic structure

has a grain size of 1, the remaining being 2.

Decausc rare earth boron steel does not contain

strong carbide formers, the amount of carbides in the

diffused layer is not very high. The above described

metallographic structure, is superior to the chrome

steel. Rejection rate due to inferior metallographic

structure of the rare earth boron steel carburized

quenched gear is very low. On the other hand, the

gear made of chrome steel after carburization heat

treating, the rejection rate sometimes is as high as

1.5% due to simply the carbide being out of spec.

d. Hardness. The surface hardness of rare earth boron

steel after carburization, quenching, and annealing

is generally around RC59 to 62 which is comparable to

chrome steel gears. We therefore can meet the

desired requirement.

Due to the higher quenchability of the rare earth

boron steel after quenching of the carburized gear

the hardness in the center part of the gear in also

higher. Take for example, the 5/3.75 gear, the

hardness in the center of a 18 chromc-mangancsctitanium gear after carburization and quenching is

around RC35-39 while that of the rare earth boron

steel gear is around RC40 and above. This resulted

in a stronger base metal for the rare earth boron

steel gear.

P32-5
c. Distortion of Gear During Carburization and

Quenching.

Based on long term production experiences, the

distortion of the teeth of gear made by rare earth

boron steel after heat treating is very close to that

of the chrome steel however, the distortion of the

keyholes of the rare earth boron steel is higher than

the chrome steel in general. But through better data

gathering and statistical analysis, we conclude that

for the same kind of gear, the fundamental trend of

distortion is basically the same for both rare earth

boron steel and chrome steel. The distortion

direction is the same, basically shrinkage. The

standard deviation of the amount of distortion is

about the same. The average amount of distortion is

usually higher in the rare earth boron steel so it is

concluded that we have a handle on the trend of

distortion in a rare earth boron steel gear.

IV. Machine Testing the Machine Application of the Gear and the

Effect of Rare Earth and Boron in the Steel

1.' Machine Testing. The third shift gear of #27 tractors

gear box is a typical example. The wall of the gear is

very thin. It has four 2.3 mm diameter oil hole and the

inner wall is cladded with copper bushing. The loan is

also applied on the gear when the tractor is in the

reverse shift. The original material for this gear was

18 chrome-mangancce-titanium steel. In order to increase

the strength, it was Jator modified to use 22 chromemangancsc-molv steel. Nov; the gear is made using 25

manganese-titanium-boron and rare earth steel.

P32-6
A machine testing of four different kinds of steels was

performed in Jiangxi Tractor Plant to compare their

relative lives. The testing specification is as follows:

the load is 58.7 kg-meter (this is in equivalent to 2.19

times the load the third gear will take) . The speed is

about 727 rpm. The following table tabulates the results

of the test.

P32-7
Results

of

Test

Rig

Damage

Condition

Broken

at

oil

tap

(same)

c

o

Q) 0) b

K- g U

id id Si

u) in a — w 3

Broken

at

oil

tap

(sane)

u

C0)

a) s: o> c: 5 W U

O 5)H «C MsU m

ca c •- w

Carbon

Diffusion

Laver

(

u)

in

CO 00

• o\ 0 •

1 o

VD 1 r-~ o\

• • o o

0.67-0.76 0.86-0.95 0.67-0.

76

0.72-0.8 0.72-0.8

ss

RC |

Center O CN 1 m in

• • OV r-l CN

r-» v

m• m i v' 1 K? 1 V 1

Hardnc:

Surface

|

m•

»-l O 1 VD VO 62 64 62

1

,

• 09

57.5

*

graphic

ture 1

Center Lov;

Carbon

M

Lov;

Carbon

M

+

Some

F

Low

Carbon

M

(same) Lov;

Carbon

I!

+

Some

F

Low

Carbon

M

+

Some

F

Hetallo

Struc

Surface

I

Fine

M

+

Some

A

Fine

M

+

Some

A

(same) (same)

<Dg 1 Cw

Fine

M

+

Large

Amt A

Tine

to

Fracture

(hr)

616 776 825 626 708

1000 8641120

in . o\

r- c; H H

Time

to

Crack

Formation

(hr)

385 357 365 «n. m vo

o> r-i so

•h n m

203 229 419 415

iu

- s> a

4J ^1 « C £ *H 18

CrilnTi 22

CrMnMo 25

MnTiB 25MnTiB U r*

P32-8

Hartensite

F

Ferrite

A

=

Austenite
In our plant, wc have clone similar tests of a third shift

gear with a load of 29.4 kg meter, a speed of 1450 rpm.

After over 15,000 hours, the gear is still in good shape

and the test is continuing.

2. Actual Application. The main rear shaft cone gear of a

#7 tractor is one of the gears that sustains the largest

load. Originally 18 chrome-manganesc-titanium steel was

used and the life of the gear is about 2000 hours. Later

on, 22 chronic-manga nese-moly steel was used and 52.5Z of

the gears lasted less than 5000 hours. After switching

to 25 mangenese-titanium-boron rare earth steel,

according to a survey in 1977, 43% of. the gears has

reached the life of 10,000 hours. In three communes in.

several provinces, the application lives for that part

are 13,040 hours, 11,420 hours, and 13,560 hours

respectively. More than half of the time, their tractors

were used in a rice paddy. In May of 1981, we made a

survey of our users who reported their 25 manganesetitanium-boron rare earth steel central transmission gear

after 16,070 hours and 10,000 hours respectively arc

still in good shape. The teeth still shine and in

perfect shape. Also, in one of the applications in

automobile, rare earth transmission gear after 200,000 km

of traveling, the face of the gear is still in good shape

and shining. Another survey made on a East-Is-Rcd 54

tractor, one gear on that tractor lasted about 4,000

hours and the wear value of the face of the gear is about

.08 to .102 mm. Another test was done on a transmission

gear that was carburized and also nitrided. The test

results show that after 8,000 hours of application, the

maximum wear is about .05 mm. On the tractors produced

by our plant namely #12, #27, #55, and #54, the main

transmission gears all lasted more than 6,000 hours.

P32-9
Since 1969 to date, we have conducted seven users survey

in conjunction with personnel of research institute and

technical school and have conducted visits to the

customers. It Is clear that from these results not only

the 25 manganese titanium boron rare earth steel has a

higher bending fatigue resistance strength, it also has a

higher bonding strength, a higher contact fatigue

strength, and a lower crack sensitivity. The results

from tost rig and on-r.i to application arc consistent.

3. The Effect of Rare Earth and Keren in the Steel.

To answer the question why 25 manganese-titanium-boron

rare earth steel gear is better than the original steel

gear, we have conducted a research in conjunction with

the Peking University of Iron and Steel Technology and

here are the results:

a. Proper addition of rare earth in the steel

apparently can have a desulfurization effect and it

can also change the property of the inclusions. The

rare earth converts the elongated manganese sulfide

into spherical rare earth sulfide (Re^S^) or

oxysulfide (Re

2

0

2

S) . The elongated inclusion of

manganese sulfide apparently lowers the transverse

ductility of the steel and also its toughness. The

coversion effect of the rare earth on the sulfides

thus eliminated the detrimental effect of elongated

shape manganese sulfide. As a result, the strength

of the 25 manganese-titaniun-boron rare earth stool

gear is increased. In the meantime, the rare earth

in the steel converts the aluminum oxide or aluminum

compounds which have higher hardness to a rare earth

aluminum compound inclusion which has a lower

hardness. Through this process, the fatigue strength

of the gear was increased.
b. The rupture strength and toughness of the carburized

layer of the 25 manganese-titaniun-boron rare earth

steel gear are raised by the rare earth addition.

c. The gear has a higher carbon concentration (.25?.) in

its inner part. The boron raises its quenchability

.

After quenching the inner strength of the gear is

subsequently raised. Therefore, in gears having

identical carburized layer thickness, the 25

manganese-titanium-boron rare earth steel has a

deeper effective hardened layer therefore a longer

application life for the .gears.

At the present, our plant is mass producing gears for

tractox*s and automobiles by using the rare earth

boron steel and these gears have received good rating

and arc warmly welcome by the users.

This report was prepared in October 1981 by the Nanchang Gear

Plant.

JHC/cas

1/7/82

P32-11
.

.

.
POTENTIAL FOR SELECTIVE HARDENING

BY INDUCTION IN CHROMIUM FREE STEELS

By

Peter A. Hassel

1

Ajax Magnethermic Corp.

"Electromagnetic Induction," "Faraday's Law," "Transformer Effect," in

electrical jargon describe the single basic principle which makes it possible

to heat any electrical conductive material in an alternating or changing

magnetic field. This is the basis for all induction heating, a process which

is very much akin to I 2

R or the resistance heating of an electric furnace,

toaster or hair dryer--el ectri c current flowing against the inherent resistance

of the metal being heated. Losses considered undesirable in transformers,

motors and other electric induction devices are accentuated to heat and melt

metals. Inductive coupling is the key to some of the very unusual ways in

which electromagnetic induction can be applied for heating. No physical

contact is required nor is any sensible heat a part of the energy transfer

system. With heat generated in the charge proper, and surface energy transfer

problems eliminated, opportunities exist to apply large amounts of heating

energy, very high energy densities, and in controlled patterns where heating

currents are the image of the heating inductors.

One should not exclude low power, low energy induction heating, because

in specific uses, it can be an economically attractive method for heating

various products, where the process dictates this kind of treatment. Any

reasonable electrical conductor can be inductively heated, but in the manufactured product world of today, metals by far predominate the electrical conduction material category. In general, metals require considerable process

heating, and this was emphasized in a study which indicated an average of

P33-1
nine separate heating operations used in the manufacture of a typical steel

part. Recent efforts to more efficiently use energy are tending to reduce

this number, and induction heating is making some noteable contribution to

this endeavor.

Induction heating as a metalworking tool has application in many basic

metalworking processes. Some of these are melting, heating for hot working,

heat treating, joining, and a host of other miscellaneous uses in which shape

or metallurgical structure does not change. The latter include, shrinking,

applying metallic and organic coatings, stress inducement, sintering and

others which utilize the unique application capabilities of electromagnetic

induction. Alloy conservation possibilities certainly exist in induction

hardening, but other induction applications also provide potential to enhance

this effort.

Metal heat treating is certainly one of the major uses for the induction

heating technique. Today, in the various metalworking industries, most

engineering and production people are exposed to some form of induction

heating, very likely for one of their own manufacturing processes. Very

often it is one of the heating methods considered for a new application or a

new facility. At this time, induction heat treating of non ferrous metals

seems to be limited to annealing, mostly continuous annealing in uniform

sections. On the other hand, induction is broad in its use for heat treating

ferrous metals. Localized and surface hardening, through hardening and

tempering, normalizing and annealing, solution annealing and stress relieving

are some of the types of heat treatment performed by the induction method.

All types of steel and cast irons are candidates, and most every type of

ferrous metal is in some place or in some way now being induction heated.

P33-2
Some shape and part configurations offer limitations to uniform heating;

and for this reason, the induction method should not be considered a cure for

all metal heating problems. It is, however, being successfully applied to

parts and metal sections too numerous to mention, in thousands of installations

all over the world. In addition, new uses are continually being developed,

particularly where one or more of its unique capabilities can be put to work.

Induction hardening in this country first became visible in the early

1920's when a motor generator frequency converter, used in melting high grade

steel, was connected to a second station and ultimately used to surface harden

mill rolls for use in the production of cold rolled steel. The tendency of

alternating currents to travel in the surface layers of a conductor was recognized as a contributor to the desired surface heating and hardening.

Coupled with what was then considered high power density, the skin effect and

tendency to surface heat produced a roll with somewhat greater depth of

hardness as compared with the same alloy steel quenched from furnace heating.

Later, in about 1932, crankshaft hardening was tried by a manufacturer

of melting furnaces, using a high frequency converter which powered one of

his small furnaces. The real value of a very pronounced skin effect at higher

frequencies was recognized as a means for producing localized surface heating.

This test also revealed the potential for high frequency inducation to produce

wear resistant surfaces on cast iron and steel parts. Moreover, this treatment could be confined to specific areas where required. Beyond the value of

a single stage, short cycle, low distortion, case hardening treatment, an

additional economic advantage was recognized by progressive metallurgical

thi nkers--el imi nati on of expensive alloy extras in the purchased steel.

P33-3
Internal hardening of engine cylinder liners was developed later to

provide long life, wear resistant surfaces for use in heavy duty engines.

Axle shafts and other overall surface hardening applications must be included

as part of induction's role in reducing the need for critical alloys. Many

tons of now critical alloys have been saved in the manufacture of parts which

are locally heat treated, in the sense of controlled depth only. High density

surface heating, or controlled depth heating with induction does much to

eliminate the need for alloys in heat treating steel; when one considers the

primary role of alloys is to improve hardenabi 1 ity. Furthermore, typical

short induction heating cycles, with zero time at temperature, have been

found to render most carbide forming alloys ineffective; at least when

conventional heat treating temperatures are used.

In the late 30' s, induction heat treating was primarily seen as an

innovative, localized surface hardening technique, which could take the place

of lengthy surface addition case hardening treatments, such as carburizing

and nitriding. Selective heating not only saved time and energy, it was also

touted for reduced distortion, and some saw the potential for use of leaner,

less expensive steels.

World War II and its concerted effort to improve ordnance components,

armor piercing shot, gun parts, tank and military vehicle engine and drive

components, all provided impetus to the increasing use of induction hardening.

Metalworking industries came out of the war with considerably more knowledge

of induction heating, however, not necessarily directed to any kind of

conservation.

P33-4
Ultimate hardness was one of the prime objectives sought in making armor

piercing shot; an extremely hard nose, tapering into a more ductile structure.

In test after test, induction heating and quenching provided hardness a few

points higher on the Rockwell "C" scale as compared with identical shot

similarly quenched after heating in a fuel fired or electric furnace.

Simultaneously, this phenomena was observed in other parts, and subsequently

dubbed "Superhardness." Conservative metallurgists called it a figment of

the induction hardeners imagination, but it was there and measureable in many

hardened parts. Later investigation indicated the added hardness was a result

of rapid heating, which engendered a condition of higher stress in the hardened

martensitic structure. Some studies indicated part of the increased hardness

was due to less retained austenite in the hardened structure, when quenched

from a rapidly produced, non-homogeneous austenite.

Analysis of the forces applied in high differential induction heating

and quenching is difficult. Involved is differential thermal expansion; a

layer of plastic metal upsetting over colder hard metal, cooling in two

directions; to the quench and to the core, and all this coupled with increased

volume in the final hardened structure. Residual stresses from induction

hardening can be significant, can be beneficial; and properly applied, they

can very materially improve the fatigue strength and life of parts subjected

to torsional and bending loads. Coupled with a higher case hardness, in

which depth exceeds the applied stresses, residual compressive stresses will

do a great deal to increase endurance limits performance capabilities of a part.

P33-5
Short cycle induction heating was difficult to understand, and early

users tended to treat this type of heating as furnace heating without regard

for time. In the late 1940's and early 1950's, with considerable interest in

induction hardening, the effects of short cycle, zero time at temperature

reactions were the subject of many metallurgical investigations. A great

deal was learned to provide a substantial base, but this kind of knowledge is

never complete, and we are still on a learning curve to fully utilize the

advantages offered by extremely fast controlled heating that can also be

applied in well defined patterns.

To gain some perspective on how induction heating can reduce the need

for alloys in steel, and particularly chromium, we know hardenabi 1 i ty improvement is the primary function of most alloys used in hardenable steels. Other

than induction or direct flame hardening, furnace hardening treatments for

steel involve surface additions to case harden, i.e., carburizing, nitriding,

carbonitriding, or bulk quenching and tempering Carbon and nitrogen surface

additions produce a hard wear resistant surface, and quench and temper offers

some wear resistance, but primarily it is a process for increasing tensile

strength with some compromise in ductility. Both are candidates for, and

both have been widely applied to alloy steels, where the primary function of

alloy is to improve hardenabi 1 ity; increasing the depth to which quenching is

effective.

The mechanics of adding carbon or nitrogen involve placing these elements

in contact with the surface of the metal; atoms are exchanged and a process

of diffusion carries the added elements into the surface layers. Carburizing,

P33-6
nitriding and carbonitriding are processes which require considerable time,

measured in hours, depending upon the depth required. For this reason, there

are economical and practical limits to the maximum depth of hardend layers

produced by carbon or nitrogen addition. An increased carburizing temperature

can be used to accelerate the diffusion rate, but the process is normally

carried on a temperatures where added time or temperature may cause undesirable

grain growth.

Case depth limitations in parts hardened by carburizing or nitriding

often dictate the need for strengthening of support metal under the case.

Hardenabi lity alloys are generally specified to insure transformation at the

required depth. If a part is to resist bending or torsional loads, this

again may call for alloys to develop core properties to an adequate depth.

Alloys will be used to improve hardenabi lity so the effect of quenching can

reach the sub case or core metal. For this reason, many carburizing and most

nit riding grades of steel are alloy steels.

Other than electron beam or laser hardening, induction heating can

provide the highest energy density for controlled depth heating. Skin effect

is ever present because it makes the induction process work. Increasing

frequency tends to promote surface layer heating, and this is another element

which can be used to selectively heat only the metal required to produce

desired results.

Controlled depth heating dramatically reduces the need for hardenabi 1 ity

enhancing alloys simply by reducing the amount of heat that must be removed in

the quenching process. In the critical early stages of quenching, heat moves

P33-7
in two directions: to the quench and to the core, even in a part surface

heated overall. For example, in a deep induction case hardened, 2" diameter

shaft with 3/16" effective hardened depth, the total heat content is less

than one half that of the same part out of a furnance at 1650°F.

When a shaft is locally heated, for example, in a bearing area, the heat

also moves axially, and the cooling effect with added adjacent colder metal

is even more dramatic. Scan or progressive induction hardening, widely used

in heat treating shafts, gears, machine tool ways and other parts, in effect

is a continuous incremental heating and cooling process where heat conduction

of quenching takes place in all directions. Critical cooling rates to

substantial depths are easily achieved for austenitized low hardenabi 1 i ty

steel s--depths which will support reasonably hgih loads in pressure or

brinelling, torsion or bending.

Alloying elements are not always effective in induction hardening cycles.

This is particularly true of carbide formers such as chromium, molybdenum,

vanadium and tungsten, which in short induction heating time (unless abused)

precludes any appreciable grain growth. In most induction heat treating

cycles, the time at grain coarsening temperatures is such that no noticeable

growth occurs. This has been demonstrated in steel where the surface has

been close to the melting point.

Abrasion resistance imparted by undissolved carbides in hypoeutectoid

steel is know, and this may be reason to use some carbide forming alloys;

however, if higher hardness is obtainable with induction heating and quenching,

the use of hardness rather than alloy carbides might be considered in a

materials conservation program.

P33-8
The other beneficial effects cited for carbide forming alloys, including

resistance to softening and loss of strength on heating and a better balance

between strength and toughness, are not fully realized in the critical hardened

metal, unless the carbides are dissolved in the austenite prior to quenching.

This thought might be useful when considering induction hardened alloy steels,

and help in the conservation of critical category alloys, of which chromium

is a prominent member.

Like most metallurgical reactions, carbide dissolution is effectively

accelerated by increasing the austenitizing temperature. Measures induction

heat treating temperatures of 150 to 300 degrees F above furnance temperatures

are commonly used to increase carbide solution in the austenite. Grain

enlargment, largely detrimental to physical properties in steel, including

notch toughness and transition temperatures previously indicated, is very

substantially elevated temperatures, the time element for any measurable

grain growth does not exist. This has been well proven in extremely short

heat cycles with measured surface temperatures will over 2000°F. The timetemperature relationship of metallurgical reactions is vividly demonstrated

in almost all induction heating applications for metal working, in particular

selective surface hardening.

In reality, where heat is localized in depth or in area, the induction

hardening process is capable of providing, without presence, the benefits of

most commonly used alloys of engineering steels. Potential benfits on a

comparable basis include increased hardenabi 1 i ty and improved physical properties to meet the application demands of products requiring resistance to

failure from abrasion, bri nel ling, torsion, bending and low temperature

P33-9
impact. The ability to develop desirable residual stresses must be recognized

because, properly applied, these can substantially improve the performance of

heavily loaded parts. In some applications, fully utilized on an educated

basis, the total metal in a part can be reduced, and alloy conserved as well.

The potential use of induction hardening to reduce critical alloy consumption within its scope is real, but the net contribution will depend upon how

seriously metalworking industries consider this a need, either for economic

reasons or survival. If the philosophy becomes pervasive, the savings can be

substantial. This has been proven by some of the more dedicated, progressive

metallurgical thinking companies who are substantial users of steel. Induction

hardening has played a major role in some long term programs for better

utilization of materials.

Many products made of chromium containing alloy steels can be relegated

to grades using other alloying elements for hardenabi 1 i ty or abrasion resistance. Manganese, molybdenum and boron are useful substitutes where alloy is

needed, but induction has turned many alloy steel parts to plain carbon

grades without sacrificing physical properties, consuming more energy or

adding to manufacturing cost. Material and heat treatment should be considered

when a part is in the development stange to determine how it can be designed

to best utilize materials and processes. In a dedicated materials conservation

philosophy, material is the first consideration, followed by an intensive

search for the best process to develop needed performance properties.

In this country, an abundance of materials and energy has made us somewhat

complacent about conservation; and now our friends from countries where

resources have been limited, are showing us how to use methods we developed

P33-10
to produce higher quality products with less critical materials. Fortunately,

we have a few outstanding examples of industries with foresight who are ahead

in the material and process technologies. They are very observant, international

in their outlook, and have a solid philosophy, developed over many years. We

can all take lessons for the future from those outstanding examples.

Briefly we have shown how some selective induction hardening can reduce

or eliminate the need for critical alloys in heat treating engineering steels.

We haven't begun to scratch the surface, and a good start is to look at every

existing induction hardened part made of alloy steel. Beyond that, no one

should write a specification calling for alloy induction hardened steel, and

particularly those containing chromium, without asking if the alloy is needed,

or more importantly, will it be useful induction hardened. Outmoded material

specifications should be examined, and standards should be opened up for

intelligent evaluation of new knowledge on processes affecting metals and

alloys. We can no longer stand extreme conservatism in government and military

specifications or in engineering materials codes.

Induction heating can undoubtedly contribute much to conserve or eliminate

chromium in applications other than hardening. Many of these possibilities

are fact and have been proven in laboratory or behind the scene uses, ready

to be applied to products in which alloys and critical materials are being wasted.

Some possibilities are in unsuspecting areas, such as an induction

treatment to impart corrosion resistance. Will we let others find the needs

of the future, and if profitable and advantageous to them, show us; or do we

have the foresight and the desire to properly use what has been generously

provided in this finite world?

P33-1I
.
1982-09-09

BEARING STEELS OF THE 52100 TYPE WITH REDUCED CHROMIUM

by C. F. Jatczak, The Timken Company

The Timken Company is primarily a manufacturer of tapered roller

bearings which are made from carburizing steels that generally

contain low amounts of chromium - typically between 0.25 and

0.50%. At the same time, however, the Company is a major supplier

of 52100 high carbon steels to various ball bearing manufacturers.

By contrast with the carburizing grades, the 52100 steel contains

1.5% Cr. This makes our job of supplying it without interruption

very vulnerable to chromium shortages.

In anticipation of a possible Cr shortage in the early 1970's

(which later did occur), we set out to apply our combination

of steel and bearing knowledge to develop a suitable low Cr alternate which would satisfy all of the technical and bearing performance requirements of standard 52100 ball bearing components.

A secondary aim was to make the new composition so economically

attractive that our customers would be willing to evaluate its

performance as soon as possible.

The resulting composition, designated TBS-9, is shown in Figure 1.

A savings of 1% Cr was achieved. Test results to date indicate

that TBS-9 offers processing characteristics and service performance which are in fact superior to those of 52100 at a lower

cost. For example, the metallurgical characteristics of the steel

are such that its annealing times are much shorter, thus lowering

our costs. We were able to pass these savings on to our customers,

which in 1971 amounted to $60/70 per ton for steel in tubular

form furnished with a spheroidized microstructure. Today this

savings is not as lucrative as it was in 1971, due largely to

a large increase in the price of Mo; nevertheless it is still

signif icant

From the Cr conservation standpoint, the 1% lower Cr content

in TBS-9 amounts to a rather modest but important savings of

20#/ton. Based on AISI data for 1981, in which approximately

91,000 tons of high carbon bearing steels were sold domestically,

the bulk savings in Cr could amount to 910 tons annually if TBS-9

were substituted across the board.

Since 1971, TBS-9 has received wide acceptance as an alternate

for 52100. About one-third of Timken's high carbon steel business

is now TBS-9 steel.

Hardenability played an important role in the development of

this grade. The steel was designed to meet the hardenability

specifications of the 52100 grade. This was easily accomplished

P34-1
2

as shown in Figure 2, because of extensive prior work conducted

by Timken on the influence of the various alloying elements on

hardenability of high carbon steels.* The process was simply

one of substitution of Mn and Mo for Cr based on the data shown

in Figure 3

.

Metallurgical and processing characteristics were also designed

to be at least equivalent. The results were equivalent forging

and tube piercing characteristics and as already noted superior

annealing behaviour, because less time is required to produce

the spheroidized structure which provides ease of machining.

The maximum as annealed hardness is BHN 207; however, the typical

range is 179/192 Brinell.

The effects of various austenitizing and tempering temperatures

on hardness of TBS-9 are illustrated in Table 1.

The ultimate criterion for a ball bearing steel is its resistance

to rolling contact fatigue. The fatigue resistance of TBS-9 was

compared with 52100 and lower carbon version 5160 in our laboratories - our fatigue machines provide a rolling contact test

that controls all variables except the material itself. Contact

geometry is very similar to a ball bearing. The test conditions

employed were: specimen, 4 by 5/8 in. diameter cylinder; load

rings, 1 in. crown radius on an 8h in. diameter; speed, 9000

rpm; load, 1680 lb.; maximum Hertz stress, 700 ksi; lubricant,

SAE 20 oil; temperature, 175 °F.

The results are shown in Figure 4. The contact resistance of

TBS-9 is superior to that of 5160 and similar to that of 52100.

Actual bearing field tests conducted by our customers since these

RCF tests have confirmed the equality of TBS-9 with 52100 with

respect to bearing life. As was noted earlier, TBS-9 has replaced

at least 1/3 of our high carbon bearing steel orders. This is

the best confirmation of its viability as a ball bearing steel.

To summarize, it is clear that a reduction in Cr in high carbon

bearing steels need not cause any consternation in the industry

at all. Simple substitution of the TBS-9 should provide uninterrupted production of ball bearing races even during drastic shortages of the element chromium.

* See "Determining Hardenability from Composition," by

Jatczak, Metal Progress, September, 1971, p. 60.

F.

P34-2
FIGURE 1

COMPARISON OF 52100 AND TBS-9 CHEMISTRY

Type C Mn Si Cr Mo

52100 .95/1.10 .25/. 45 .20/. 35 1.30/1.60

TBS-9 .89/1.01 .50/. 80 .20/. 35 .40/. 60 .08/. 15

Table I — Effects of Austenitizing and Tempering

Temperatures on Hardness of TBS-9

Temperature, F Hardness, Rc

Austenitizing As Quenched After 300 F Temper

1,475 66.0 63.5 1,500 65.0 64.0 1,525 66.0 63.5 1,550 66.0 63.5 1,575 65.0 63.0

Tempering After 1 Hr After 2 Hr (1,550 F Quench) Temper Temper

200 65.0 65.0 300 64.0 63.5 400 60.5 60.0 500 58.5 58.0 600 55.5 55.5 700 52.5 52.0 800 49.0 49.0 900 45.0 45.0

1,000 40.0 39.0 1,100 36.0 35.0 1,200 32.0 30.0 1,300 26.0 24.0

Typical Rockwell C hardness (normalize 1,650 F, anneal, quench 1,550 F). J1 position (1/16 in. Jominy

distance), 65; J2, 65; J3, 65; J4, 61; J5, 47; J6, 43; \ 43; J8, 43; J10, 43; J12, 41.

P34-3
ROCKWELL

C

HARDNESS

SCALE

TYPE HEAT NO. C Mn Si Cr Mo

NORM.

TEMP.

°F

QUENCH

TEMP.

° F

52100 Avg 16 hts. .98/1.10 .25/. 45 .20/. 35 1.30/1.60 1650 1500

TBS-9 Avg 35 hts. .89/1.01 .50/. 80 .15/. 35 .40/. 60 .08/. 15 1650 1500

DISTANCE FROM QUENCHED END OF SPECIMEN IN SIXTEENTHS OF INCH

Figure 2. Comparison of Hardenability of TBS-9

versus 52100.

P34-4
Figure 3. Multiplying factors for the calculation of

hardenability in high carbon quenched from

1525 °F (830C), see reference indicated below.

tao

C/O

c/>

<D

O

CD

OO

Figure 4. Comparison of rolling contact fatigue life.

P34-5

I. STRUCTURAL ALLOYS

:

A. Bar Products

To begin I would like to first briefly review the state-of-the-art in con- servation and substitution technology for structural alloy bars. This

subject has been one of deep concern and some considerable research at

Republic Steel as we are the largest bar producer in the U. S. In a typical business year an analysis of Republic's consumption of raw materials

by product line reveals that approximately 80% of our purchased chromium

is consumed in the production of bar products. Our stainless operations

utilize 11% of the purchased chromium and the remaining 9% is divided

evenly between tubular and flat rolled products.

As a result of this consumption distribution of chromium we have focused

our alloy substitution research on bar products and have emphasized re- search to utilize substitutes with wide distribution and domestic sources

such as vanadium and molybdenum.

Alloy Substitution — The Current State-Of-The-Art

Alloy substitution in steels is not a new endeavor and in recent history

dates back to World War II when 16 National Emergency Steels (N.E. Steels) were developed and introduced in 1944. As the nation was faced with

shortages of critical alloying elements and of uncontaminated alloy scrap,

the steel industry united together and developed in a relatively short

period of time the N.E. Steels. The N.E. Steels were all fine-grain alloy

steels produced by the best alloy practice known to the industry and they

were developed on the basis of equaling the hardenabi 1 ity of the grades

they were to replace.

The N.E. Steels were triple alloy steels (Cr-Ni-Mo steels, 8600 and 9400

series) which permitted greater recovery and utilization of alloys from

scrap. They therefore aided in meeting mill production schedules as well

as producing steels conforming to the specified chemical ranges which was

a problem with the single or dual-alloy type steels such as 5100, 3100,

4100, and 4600. Many of these N.E. Steels such as 8620, 8630, and 8720

have been and are still the standards and "Work-horses" of industry.

The EX Steels

More recently there has been another example of large-scale alloy substitution as a result of raw material shortages. This example resulted from

the nickel mining strike by INC0 and Falconbridge Nickel Mines, Ltd.,

which occurred in 1969. The strike at the INC0 mines lasted four months

beginning on July 9, 1969, and ending on November 14, 1969. The resulting

shortage in nickel manifested itself by an intense level of development

activity on the part of two domestic steel companies and two ferroalloy

suppliers to provide alternate steels with equivalent hardenabi 1 ity and

mechanical properties with reduced nickel levels.

P35-1
- 2 -

This development activity resulted in both Republic and International Harvester introducing new economical alternative steels (the SAE EX steels)

to the marketplace by January 1970. Therefore, in about five months' time

alternate materials had been developed in response to the raw materials

shortage. By 1971, 52 EX steels had been developed and offered to the

trade. Ten years later (1981) 30 EX steels remained in existance.

The significant point to be made is the development response time required

to introduce substitute alloy grades as a result of a raw materials shortage. In the case of the nickel shortage in 1969 the steel industry and

ferro alloy suppliers had alternate grades of alloy bar products on the

market in five months time.

Today we have accurate computerized hardenability prediction systems to

aid us in our alloy development work which enables us to shorten the development cycle of alloy bar products even further. The development of new grades of stainless steel is a different story and takes a considerably

longer period of time to develop (usually several years).

Current Alloy Substitution Technology

In this section a brief description of existing alloy substitution technology will be presented. This information is intended to illustrate the

current technolgoy and know-how that we have at our fingertips which allows us to respond to customer demands for alloy substitution or raw materials shortages as they arise.

Alloy Substitution Capabilities Within Current AISI-SAE Grades

Illustrated below are a few examples of where alloy substitution can readily take place by merely substituting one existing AISI-SAE grade for an- other based on equivalent hardenability bands. The first example illustrates how 8600 series steels (Ni-Cr-Mo grades) can be substituted for the

5100 series steels (straight Cr grades) and reduce the average Cr content

from about 0.80% to about 0.50%. The second example deomonstrates how a C-Mn-Mo grade (4042) can be substitued for a C-Mn-Cr grade (5046) in the

event of a Cr shortage. The third illustration shows how a C-Mn-B steel

(15B37) can be substituted for a C-Mn-Cr steel (5132) in the event of a Cr

shortage.

P35-2
Rockwell

Hardne

3

C-Mn-Cr Ni-Cr-Mo

C-Mn-Cr Ni-Cr-Mo

C-Mn-B C-Mn-Cr

Distance From Quenched End of Specimen, Sixteenths

These few examples should suffice to illustrate how existing standard AISI

grades of steel can be substituted for other existing grades in the event

of a raw materials shortage.

P35-3
- 4 -

Alloy Substitution Capabilities Via the SAE EX Steels

As mentioned previously, the SAE series of EX steels represent the SAE

numbering system for new grades of alloy steels. As a function of time

the EX steel either is dropped because of lack of interest in the alloy or

it becomes a standard SAE grade if the alloy becomes popular and widely

used. The currently available EX steels and the standard SAE grade for

which they are normally substituted are shown in the list.

THE EX STEELS AND EQUIVALENT STANDARD GRADES

Composition (%)

EX Equivalent

No. C Mn Cr Mo Other SAE Grade

10 0.19-0.24 0.95-1.25 0.25-0.40

15 0.18-0.23 0.90-1 .20" 0.40-0.60

16 0.20-0.25 0.90-1.20 0.40-0.60

17 0.23-0.28 0.90-1 .20 0.40-0.60

18 0.25-0.30 0.90-1.20 0.40-0.60

19 0.18-0.23 0.90-1 .20 0.40-0.60

20 0.13-0.18 0.90-1.20 0.40-0.60

21 0.15-0.20 0.90-1.20 0.40-0.60

24 0.18-0.23 0.75-1 .00 0.45-0.65

30 0.13-0.18 0.70-0.90 0.45-0.65

31 0.15-0.20 0.70-0.90 0.45-0.65

32 0.18-0.23 0.70-0.90 0.45-0.65

33 0.17-0.24 0.85-1.25 0.20 min

34 0.28-0.33 0.90-1 .20 0.40-0.60

36 0.38-0.43 0.90-1.20 0.45-0.65

38 0.43-0.48 0.90-1 .20 0.45-0.65

39 0.48-0.53 0.90-1.20 0.45-0.65

40 0.51-0.59 0.90-1 .20 0.45-0.65

54 0.19-0.25 0.70-1.05 0.40-0.70

55 0.15-0.20 0.70-1 .00 0.45-0.65

56 0.08-0.13 0.70-1.00 0.45-0.65

57 0.08 max 1 .25 max 17-19

58 0.16-0.21 1 .00-1.30 0.45-0.65

59 0.18-0.23 1 .00-1 .30 0.70-0.90

60 0.20-0.25 1.00-1.30 0.70-0.90

61 0.23-0.28 1.00-1.30 0.70-0.90

62 0.25-0.30 1.00-1.30 0.70-0.90

63 0.31-0.38 0.75-1.10 0.45-0.65

64 0.16-0.21 1.00-1.30 0.70-0.90

65 0.21-0.26 1.00-1 .30 0.70-0.90

0.05-0.10 0.20-0.40 Ni 8620

0.13-0.20 - 8620

0.13-0.20 - 8622

0.13-0.20 - 8625

0.13-0.20 - 8627

0.08-0.15 0.0005 B min 94B17

0.13-0.20 i 8615

0.13-0.20 - 8617

0.20-0.30 - 8620

0.45-0.60 0.70-1 .00 Ni 4815

0.45-0.60 0.70-1.00 Ni 4817

0.45-0.60 0.70-1 .00 Ni 4820

0.05 min 0.20 Ni min 4027

0.13-0.20 - 8630

0.13-0.20 - 8640

0.13-0.20 - 8645

0.13-0.20 s 8650

0.13-0.20 - 8655

0.05 min _ 4118

0.65-0.80 1 .65-2.00 Ni 4817

0.65-0.80 1 .65-2.00 Ni 9310

1 .75-2.25 - .. 30303

_ 41 18

- - 8620

_ _ 8622

- - 8625

- - 8627

0.0005-0.003 B

P35-4
- 5 -

Currently there are about 30 different EX steels which can be substituted

for a large variety of standard SAE grades. Generally the use of an EX

steel results in a reduction in either the Ni content or the Mo content

compared to the standard SAE grade. As we have experienced both Ni and Mo

shortages in the past 10 to 12 years the utilization of the SAE EX steels

has been both industrially and technologically significant.

Alloys "On the Shelf" as a Result of Recent Research

Recent Republic Research efforts to develop bar products with reduced

levels of critical alloying elements while maintaining the required heat

treatment response and mechanical properties have resulted in a number of

"on the shelf" alloys ready to go as demand arises. The table below sum- marizes some of these developments.

NEWLY DEVELOPED PRODUCTS WHICH ARE EXAMPLES OF MATERIALS SUBSTITUTION

Product Area Old Grade New Grade Alloy Savings

A1 loy Bar 3310 3120 Reduce Ni from 3.50 to 1.75%

Modified Reduce Cr from 1.50 to 0.60%

A1 loy Bar 4815 3120 Reduce Ni from 3.50 to 1.75%

Modified Reduce Mo from 0.25 to 0%

A1 loy Bar 4620 3120 Reduce Mo from 0.50 to 0%

Modified Modified

A1 loy Bar 4815 "2Ni -IMo" Reduce Ni from 3.50 to 1.75%

A1 1 oy Bar EX-55 "2Ni -IMo" Reduce Cr from 0.55 to 0%

A1 loy Bar 4118 1 OBI 8 + V Reduce Mo from 0.11 to 0%

Reduce Cr from 0.50 to 0%

A1 1 oy Bar 4027 C-Mn-Cr Reduce Mo from 0.25 to 0%

U1 tra--High- HP 9Ni -4Co "lONi-lMo" Reduce Co from 4.5 to 0%

Strength

A1 loy Steel

The first alloy shown (3120 Modified) is actually now off the shelf and in use in automotive steering gear systems. A second alloy "2 Ni - 2 Mo" is

currently being evaluated for use in rock-drilling bits and would reduce

Cr content from 0.55% to 0% compared to the currently used EX-55 steel.

In the vacuum-melted specialty steel area "10 Ni - 1 Mo" steel is under

development as a replacement for the HP 9 Ni - 4 Co steels in the event of

a cobalt shortage.

P35-5
- 6 -

Computer Prediction Capability as an Alloy Development Tool

In the past several years various hardenabi 1 ity prediction models have

been developed. The most accurate computer prediction model is the Minitech hardenabi 1 ity and mechanical property predictor. This prediction

system provides a rapid and highly accurate means of predicting hardenabi 1 ity, heat treatment response, and mechanical properties as a function

of steel composition. This therefore provides a rapid and reliable means

of designing alternate bar product alloys based on current alloy costs and

availability without going through a lengthy program of melting and evaluating experimental steel compositions.

Examples of the type of information that the Minitech predictor can provide and the accuracy of the calculations are shown below. The first and

most important information desired is the core hardenabi 1 ity of the steel.

Shown in Figure 1 are comparisons of the experimental and Minitech predicted hardenabi 1 ity curves for a given composition of 8620 steel. The

excellent agreement between experimental and predicted results in apparent.

For carburizing grade steels such as 8620 the next piece of information an alloy developer needs is the case hardenabi 1 ity or carburizing response of

the alloy. Shown below in Figure 2 is an example of the ability of the

Minitech system to calculate or predict the case hardenability for a given

set of heat treatment conditions and steel composition. Again, excellent

agreement is observed between experimental and predicted results.

P35-6
Hardness

(Rq)

- 7

-

70

60

50

40 h

30

20

10

Grade ~~

C

0.21

!

?: 8620

Mn P S 3i Ni Cr Mo Cu A1

-

0.83 0.008 0.020 0 31 0.52 0 .51 0.16 3.17 0.038

O Experii

Predic

nental

ted (Min it

i

i

I

i

i

[

A ech)

<r £

!

<0

I.

-

1

S t

)( &

6 a

1 Qa

$

l

8 12 16 20

Distance From Quenched End (1/16")

24 28

Figure 1. Comparisons of the Experimental and Minitech Predicted

Hardenabi 1 i ty Behavior for 8620 Steel.

P35-7
Hardness

(Rq)

- 8 -

70

60

50

40

30

20

10

i

!

1 1

i

!

i

; 1

O Experimental

A Predicted (Minitech)

: i

.

i

i

1

:

i a s 9 o L A (A

A

S

i 3 j $ A

c^ A ' 0 A

Grad — C

! 1

e: 8620

Mn P S Si Ni Cr Mo Cu Al

o. is o .82 o.o;

Carbon Potent'

Carburized 17(

Held at 1575

13 0.021 0

'al = 0.9C

30 F - 2.5

r - 0.5 Hr

.31 0.44

%

Hrs, Coo'

, Oil Quer

3.57 0.1 7

led in 1 H ich, Tempe

0.09 0.03

r to 1575

r 300 F

- F,

1 . 5 Hrs

0.010 0.020 0.030 0.040

Case Depth (inch)

0.050 0.060 0.070

Figure 2. Comparisons of Experimental and Predicted Case Hardenability Response for 8620 Steel.

P35-8
- 9 -

The next information that is desired, for a new alloy, are the mechanical properties of the heat-treated steel for various section sizes of

bar products. Shown below are the calculated (Minitech) and experimental mechanical properties of pseudo-carburized (core properties)

8620 steel

MECHANICAL PROPERTIES (EXPERIMENTAL VERSUS PREDICTED)

OF PSEUDO-CARBURIZED 8620 STEEL

Bar Diameter

(inch)

(1700

BHN

-v 1550 F,

UTS (k si

)

Oil Quench,

YS (k si

)

300 F Temper)

% Elonq % RA Method

1/2 388 199,500 157,000 13 49 Experimental

399 199,000 181,000 14 47 Predicted

1 255 126,700 83,750 21 53 Experimental

268 134,000 112,000 21 60 Predicted

2 235 117,200 73,000 23 58 Experimental

235 118,000 96,000 24 64 Predicted

Considering the normaT heat-to-heat variability this agreement between calculated and experimental mechanical properties is considered

to be good.

For through hardening steels such as 4140, the quenched and tempered

mechanical properties as a function of tempering temperature are of

interest to the design engineer and the metallurgist. A comparison

of the quenched and tempered mechanical properties (experimental versus predicted) are shown in the table below for two different grades

of alloy steel

QUENCHED AND TEMPERED MECHANICAL PROPERTIES

(EXPERIMENTAL VERSUS PREDICTED) FOR 4140 AND 8650 STEELS

Temper RA

Temperature

Grade (F)

YS (k si

)

UTS (ksi) % Elong %

Exp Pred Exp Pred Exp Pred Exp Pred

4140 400 238 235 257 247 8 11 38 38

600 208 199 225 215 9 13 43 44

800 165 169 181 188 13 15 49 49

1000 121 128 138 149 18 19 58 57

8650 400 243 268 281 275 10 10 38 32

600 225 221 250 235 10 12 40 41

800 192 196 210 213 12 13 45 45

1000 153 146 170 166 15 17 51 54

1200 120 113 140 135 20 21 58 60

P35-9
10 -

These data, generated for 1-inch round bar product, demonstrate good

agreement between experimental and predicted values. In addition to the

value of this powerful prediction tool to the alloy developer, the production metallurgist can use this computer program very successfully as a quality control tool and in heat disposition.

In summary, there exists today a very large body of knowledge concerning

the metallurgy and application of carbon and alloy bar products. This

knowledge bank includes, in part, the development, application and over 35

years of experience with the WWII National Emergency Steels, the SAE EX

steels most of which originated from the 1969-1970 time period when there

was a nickel shortage due to a strike, and for the last several years includes computerized hardenabi 1 ity prediction systems which are highly ac- curate and extremely fast and result in a significant shortening of the

alloy development cycle time. With this knowledge base the steel industry

is positioned to respond very rapidly to any possible raw materials shortage.

In terms of alloy bar steels there would not be a serious problem in the

event of a chromium shortage. The technology is in hand to readily find

alloy substitutes for Cr containing alloy steels through the judicious use

of Mo, B, and V. Whether it is a carburizing grade such as 8620 or a through hardening grade such as 4140 there is not a problem in finding an adequate non-chromium containing substitute.

B. Structural Alloy Plate

Under the category of structural alloy plate materials we will briefly

discuss the possibilities for chromium conservation and alloy substitution

in the following classifications of plate materials.

1) Constructional Alloys

2) Abrasion-Resistant Alloys

3) Weathering Steels

4) Pressure Vessel Steels

1 ) Constructional Alloys

As shown in the table below, many of the high strength 80 to 100 ksi

Y.S. structural, weldable plate steels contain Cr and Mo. These

quenched and tempered plate materials (covered under ASTM Specifications No. A514 and A517) are produced in a variety of thicknesses up

to 6" and are intended for use in welded bridges and other structures.

P35-10
11

Constructional Alloys (High-Strength Plate)

Composition

Name C Mn Si Ni Cr Mo Other

Y.S.

ksi

ASTM

Spec

|| SSS-100 0.16 0.55 0.27 1.70 0.50 Cu, V,B 100 A514

RQ100 A 0.16 0.60 0.27 - - 0.55 B 100 A51

N-A-XTRA-100 0.21 0.85 0.65 - 0.65 0.28 Zr,B 100 A514

T-l 0.15 0.80 0.20 0.85 0.55 0.55 Cu,B,

100 A514

HY80 0.18 0.25 0.27 2.75 1.40 0.35 - 80 - HY100 0.20 0.25 0.27 3.00 1.40 0.35 - 100 - HY130 0.12 0.75 0.27 5.00 0.55 0.45 V 130 -

In ASTM Specification A514 there are fifteen different compositions or Grades listed (Grades A through Q) some of which are shown in the

table above. Of these fifteen different grades there are four grades

that do not contain chromium. One of these grades is shown in the

table above and it can be seen that higher Mo levels and the addition

of boron are utilized in order to meet the hardenabi 1 ity requirements. The different compositions indicate that the balance of alloying elements is primarily based on hardenabi 1 ity considerations and

experience has indicated that the non-chromium containing grades meet

the specification and customer requirements equally as well as the Crcontaining grades. Similar to the situation with heat treated bar

products it is preimarily a matter of achieving hardenabi 1 ity most

economical ly.

Therefore, if there were a chromium shortage there would not be a technical problem in the area of quenched and tempered constructional

plate steels.

The situation with HY80, HY100, and HY130 steels for Naval ship applications is not as straight forward and adequate substitutes for these

grades are not readily apparent at this time. Later in this session

Dr. Brian Jones will have information to share concerning the possibility of substituting microalloyed HSLA plate materials for HY80

steel

2) Abrasion-Resistant Plate Steels

A similar analysis of abarasion resistant plate steels shown by the

compositions listed in the table below reveals that there is an ample

choice of quenched and tempered (Q & T) abrasion resistant plate materials that are Cr free. The users can select from a variety of C-Mn,

C-Mn-Mo-B, and C-Mn-Mo-Ni-B Q & T alloys. Therefore, chromium availability would not be a problem for abrasion resistant plate steels.

P35-11
12 -

TABLE II

ABRASION-RESISTANT ALLOYS

Composition, %

Name C Mn P S Si Cu Cr Mo Other fa)

Hardness,

Bhn

SSS-AR-321 0.25 0.40-0.70 0.04 0.05 0.20-0.35 0.20-0.40 0.85-2.0 0.15-0.60 Ti or V, B 321

SSS-AR-360 0.25 0.40-0.70 0.04 0.05 0.20-0.35 0.20-0.40 0.85-2.0 0.15-0.60 Ti or V, B 360

SSS-AR-400 0.25 0.40-0.70 0.04 0.05 0.20-0.35 0.20-0.40 0.85-2.0 0.15-0.60 Ti or V, B 400

RQAR-321 0.25-0.32 0.40-0.65 0.035 0.040 0.20-0.35 0.80-1.15 0.15-0.25 _ 321 RQAR-340 0.25-0.32 0.40-0.65 0.035 0.040 0.20-0.35 - 0.80-1.15 0.15-0.25 - 340 RQAR-360 0.25-0.32 0.40-0.65 0.035 0.040 0.20-0.35 - 0.80-1.15 0.15-0.25 - 360

RQAR-400 0.25-0.32 0.40-0.65 0.035 0.040 0.20-0.35 - 0.80-1.15 0.15-0.25 - 400 RQC-321 0.28 1.50 0.040 0.050 0.20-0.60 - - - B 321 RQC-340 0.28 1.50 0.040 0.050 0.20-0.60 - - - B 340 RQ321A 0.12-0.21 0.45-0.70 0.035 0.04 0.20-0.35 - - 0.50-0.65 B 321 RQ340A 0.12-0.21 0.45-0.70 0.035 0.04 0.20-0.35 - - 0.50-0.65 8 340 RQ360A 0.12-0.21 0.45-0.70 0.035 0.04 0.20-0.35 - - 0.50-0.65 B 360

RQ321B 0.12-0.21 0.45-0.70 0.035 0.04 0.20-0.35 - - 0.45-0.60 1.20-1.50 Ni, B 321

RQ340B 0.12-0.21 0.45-0.70 0.035 0.04 0.20-0.35 - ; - 0.45-0.60 1.20-1.50 Ni, B 340

RQ360B 0.12-0.21 0.45-0.70 0.035 0.04 0.20-0.35 - - 0.45-0.60 1.20-1.50 Ni, B 360

3) Weathering Steels

Weathering steels are HSLA 50 ksi min Y.S. plates and shapes intended

primarily for use in welded bridges and building were weight savings

and durability are important. The ASTM A-588 standard specification

covers ten different grades of weathering steels (Grades A through K)

and three of these grades do not contain chromium. While most of

these steels contain C-Mn-Si-Ni-Cr-Cu and V they can be made to have

adequate atmospheric corrosion resistance, in the absence of Cr, by

increasing the Ni and Cu levels. A second option would of course be

to do without weathering steels and get out the bucket of paint and

paint brush.

The bottom line with weathering steels is that an absence or shortage

of Cr would not be a problem.

4) Pressure Vessel Plate Steels

ASTM Standard Specification A387 for pressure vessel plates, covers

the Cr-Mo alloy steel plates intended .primarily for boilers and pressure vessels designed for elevated temperature service. This specification covers the well known and widely used 2^1/4 Cr - 1 Mo and 1-1/4

Cr - 1/2 Mo steels.

\ \

P35-12 \

\\
13 -

The market for the ASTM A-387 grades (there are 8 different grades in

this specification) is some 8,000 to 9,000 tons/year domestically and

15,000 to 20,000 tons/year worldwide. Concerning plate products there

is only one major domestic producer (Lukens Steel) the remainder of

the steel being foreign, primarily from Japan. The most popular grade

is Grade 11 (1-1/4 Cr - 1/2 Mo) which is used extensively in process

piping and heat exchanger tubing. The next most popular grade is

Grade 22 (2-1/4 Cr - 1 Mo) which is primarily used as a plate product

for pressure vessels and high temperature, high pressure hydrogen ser- vice for petrochemical applications.

In these grades if you lower Cr content you loose high temperature

strength. As the writer understands the state of the art of technology in this class of steels there are no known substitutes for these

grades and furthermore there is apparently no research work being done

to find substitutes or reduce Cr levels in these pressure vessel

steels. The current thinking is that there is no known substitute for

Cr when it comes to resistance to hydrogen at elevated temperatures.

The Climax Molybdenum Company is doing research to develop an alloy

which will have greater high temperature strength, greater hardenability and greater resistance to hydrogen. This has lead to the

development of a 3 Cr - 1-1/2 Mo - .10 V steel. Unless the section

sizes can be reduced significantly, this development, while technologically significant, will not be beneficial from a chromium conservation viewpoint.

In this product area the bottom line would appear to be that Cr substitution or conservation does not look favorable at this time and

there is a real need for research in this area. As there are a limited number of domestic suppliers of A-387 plate steels this is an area where the Federal Government should step in and sponsor some R & D work.

II. HIGH STRENGTH LOW ALLOY STEELS

The high strength low alloy (HSLA) steels and their possible vulnerability to

a chromium shortage will be reviewed in depth by the next two speakers on the

program; Mr. Michael Korchynsky of Union Carbide Corporation and Dr. Brian

Jones of Niobium Products Company. Their presentations will cover plate

steels, skelp for line pipe applications, dual-phase steels HSLA microalloyed

steels, weathering steels and HSLA steels in general. Dr. Koo from Exxon Research will discuss the optimization of alloying elements in low carbon dual

phase steels and medium carbon structural steels.

III. ULTRAHIGH STRENGTH STEELS

The major use of ultrahigh strength specialty steels is in aircraft forging

applications. These critical load bearing forgings are used in such applications as engine mounts, tail section forgings, wing mounts, flap tracks and

landing gears to mention just a few. The landing gear typically comprises

about 14% of the weight of the aircraft, therefore, high strengths and good

fatigue resistance are mandatory.

P35-13
- 14 -

The primary ultrahigh strength steels which have been used for landing gears

over the past 30 to 40 years are shown in the table below. What is readily

apparent from this list of steel compositions is that all of the alloys contain chromium. What is not nearly as readily apparent is whether a chromiumfree steel could be developed for this very demanding high hardenability

appl ication.

Aircraft Landing Gear Steels

Composition

A1 loy C Mn Si Ni Cr Mo V Other

98 BV 40 0 . 43 0.85 0.65 0.75 0.90 0.50 0.04 B

4330 M 0 .,30 0.95 0.27 1.80 0.85 0.40 0.08 - 4340 0 . 40 0.75 0.27 1.80 0.80 0.25 - - 300M 0 .,40 0.75 1.65 1.80 0.80 0.40 0.80 - HP 9-4-30 0 . 30 0.25 0.10 7.50 1.00 1.00 0.10 4.5 Co

HP 310 0 ..40 0.75 2.50 1.80 0.90 0.40 0.20 -

There is not currently any alloy development work in progress which has as its

goal to develop a Cr-free steel for ultrahigh strength steel aircraft landing

gear applications. As both military and commercial aircraft utilize the same alloys for landing gears (300M being the current work horse alloy for landing

gears) this would appear to be another very appropriate and vital area for the

Federal Government to sponsor alloy substitution research.

SUMMARY

In summary, the state-of-the-art of substitution for chromium in structural

alloy, HSLA, and ultrahigh strength steels has been reviewed. With regard to

heat treatable structural alloy bar products there are ample chromium-free

substitutes for most grades and a high technology base for rapidly developing

alternate chromium-free structural alloy bar steels. Concerning both structural alloy plate steels and abrasion-resistant plate steels there has been

available, for a considerable period of time, adequate Cr-free steels in both

of these categories of Q & T plate materials. Likewise for ASTM A-588 weathering steels, there are a number of existing Cr-free weathering plate steels.

The as-hot rolled microalloyed HSLA plate, strip and skelp steels generally do

not contain chromium and these steels would not appear to be a problem area in

the event of a chromium shortage.

The two major alloy steel areas which would be severely crippled in the event

of a chromium shortage, in the near future, are the ASTM A-387 pressure vessel

plate steels and aircraft forging and landing gear steels. For both of these

classes of steels there are no known Cr-free substitute alloys and there is no research in progress to develop such substitutes for either the pressure vessel steels or landing gear steels. Both types of steel are utilized in critically important applications and it is felt that these two areas should re- ceive serious consideration by the Federal Government to sponsor research work

to develop adequate chromium-free substitute materials.

P35-14
Alternative Compositions for Future HSLA Steels

The technology of high strength low-alloy (HSLA) steels has made

tremendous advances over the past decade. This has come about because of

the development of controlled-rolling techniques and research into the

benefits of microalloys. There now exists a family of micro-alloyed

steels capable of providing high yield strength combined with low temperature toughness and weldability. They have been utilized primarily in the

pipeline and automobile industries. The purpose of this paper is to

illustrate the properties available from this class of steel and to

suggest that their field of application be extended to include areas

traditionally reserved for heat-treatable chromium-containing steels.

The Development of HY 80

In the mid 1950's, design requirements called for a steel with 80 ksi

yield strength for use as plates and frames in submarine pressure vessels.

This necessitated the steel to provide adequate strength, ductility and

notch toughness together with good weldability under shipyard conditions.

Strength and notch toughness were, of course, also required from the weld

deposit and HAZ. In the USA, following a fairly exhaustive investigation

of the alloy approaches then available, this led to the development of the

heat-treatable Ni-Cr-Mo steels known as HY 80. These are fully-killed

aluminum-grain-refined steels which develop the desired strength properties following water-quenching from 1650°F and tempering in the range

IU 1200-1250°F. A specification for chemistry and mechanical properties for

HY 80 is given in Table 1.

P36-1
The difficulty in alloy selection was in obtaining an approach which

provided the high strength levels required together with acceptable low

temperature toughness levels. To achieve this 25 years ago it was felt

necessary to use heat-treatable nickel-containing steels which have a

shallow-sloping brittle-ductile transition curve and therefore, retain

respectable toughness at low temperatures. This however, in the steels

concerned is achieved only at the expense of weldability. Due to what is

now considered to be a relatively high carbon level and the chromium and

molybdenum additions, HY 80 is a steel which can be very prone to hydrogencracking. Its carbon equivalent using the International Institute of

Welding formula (see Table 2) is around 0.70. This means that successful

welding requires the use of low-hydrogen techniques usually in conjunction

with pre-heat to allow hydrogen escape during post-weld cooling. Three

pre-heat calculation analyses are indicated in Table 2 all recommending

levels of 150°C or 300°F minimum.

In the 1980’s, however, we no longer need to use heat-treatable nickel

steels to achieve good low temperature properties. Controlled rolling and

microalloying have now produced materials sufficiently fine-grained to

give the required strength levels and to push transition temperatures to

extremely low temperatures. Since the pipeline industry for which these

steels were primarily designed views field weldability as a primary

requisite, there has been a great incentive in this direction and these

properties are achieved with carbon equivalents below 0.35 indicating

excellent material tolerance of hydrogen and resistance to HAZ cracking.

Low Transition-Temperature Micro-alloyed Steels

Table 3 shows typical compositions of several low-carbon microalloyed

P36-2
steels which have been developed in recent years. Selection of microalloy

approach is made on the basis of final property requirements, controlled0) rolling schedules and mill capabilities.

The versatility of microalloy approach can be best illustrated by

considering the development of X70 grade (70 ksi yield strength) linepipe

steels. A few years ago it was considered necessary in producing these

materials at thicknesses of about 0.6 inches to use a Nb-Mo approach with

about 0.04% Nb and 0.30% Mo. This led in Italy for instance to the

C3)

development of the Molytar Steels. In 1979, however, due to a temporary

sharp rise in the molybdenum price, alternatives to the Mo addition were

sought. In Italy, it was found expedient at that time to use Nb-Cr steels

with 0.2% Cr. While this does not have the hardenability effect of Mo,

suitable controlled rolling schedules were easily able to achieve X70

grade requirements. Further work, however, established that even the

chromium addition could be replaced and that, with the correct heattreatment and rolling methods, a dual micro-alloy addition of Nb and V

could do the job.

The Northern Border Pipeline, therefore, which represented an order of

more than 500,000 tons of mainline 0.6 inch wall X70 grade pipe, used no

molybdenum^ and more than 90% of it, from manufacturers across the world,

used 0.04% Nb and 0.05-0.07% V as the microalloys. Some heats did utilize

the Nb-Cr approach, however, and this allowed a direct comparison to be

made of the tensile properties achieved by replacing 0.20% Cr by a much

smaller vanadium addition. It can be seen from the distribution of values

shown in Fig. 1 that no variation in yield strength attained was observed

and only minor changes in ultimate tensile strength.

P36-3
For plate of this gage and strength, therefore, there is no requirement

to use chromium and a selection of microalloy approaches are available.

For thicker grades, it is common to use molybdenum additions once again

and for high toughness at severely low temperatures - the most stringent

requirement of HY 80 specif icaiton - several novel microalloy developments

are being promoted.

Tables 4-6 illustrate a development by Nippon Steel known as ultra

(4)

low carbon bainitic steel or ULCB. Here a very small addition of boron

(.001%) leads to a transformation from ferrite-pearlite to a bainitic

structure. In order to utilize the effects of boron, nitrogen has to be

fixed by a titanium addition and carbon reduced to very low levels (< .03%).

In ultra-low carbon steels the role of niobium is also important since it

can remain in solution additionally promoting the achievement of a superfine bainite structure.

The properties achieved in this way are most impressive. It can be

seen in Table 5 that yield strengths in the vicinity of 80 ksi can be

attained in plates up to 22 mm thick with very low drop weight tear test

transition temperatures. The more commonly specified 50% Charpy crystallinity temperature is in all cases less than -80°C.

For the alloy designated UB-5 it can be seen that a yield strength of

95 ksi is achieved due to the additions of Mo and Ni , in 20 ram plate.

This excessive yield strength can in fact be traded off to allow X80

properties in plate greater than 1 inch in thickness.

Table 6 shows the obtained submerged-arc weld metal and HAZ properties. The high toughness is explained by the grain growth suppression

effect of titanium nitride particles and the elimination of martensite

formation due to te low levels of carbon.

P36-4
The requirement for very low-temperature finish-rolling (controlledrolling) can also be relaxed under certain circumstances. Recent research

(£) investigations on an HSLA steel containing 0.5% Ni and 0.13% Nb has

indicated properties in excess of X70 grade together with very good low

temperature toughness in both base plate and submerged-arc weld. These

were obtained in plate rolled to 0.6 inch thickness at a finishing temperature of 1400°F. Composition and properties of this steel are shown in

Table 7 . The low temperature toughness characteristics are illustrated

in Fig. Z .

Conclusions

The microalloyed controlled-rolled steels developed mainly for linepipe applications are able to develop 70 ksi and in certain instances, 80

ksi yield strength together with excellent low temperature toughness

properties, both in the plate material and in submerged-arc weld metal and

HAZ. They can be used in thicknesses certainly up to 1 inch and, with

additions of molybdenum, to greater thicknesses. They need contain no

chromium and offer very significant advantages in terms of weldability

over the heat-treatable Ni Cr Mo HY 80 steel traditionally used for

shipplate applications.

The microalloys involved are Nb, V, Mo and Ti. While niobium is

currently obtained largely from Brazil and Canada, there are ample deposits

in the USA, currently not regarded as commercially exploitable. The other

microalloy elements are widely distributed thoughout the world including

sources in the USA.

P36-5
References

1. "The Metallurgy and Welding of QT 35 and HY 80 Steels", The Welding

Institute. Research Report. 1974.

2. B. L. Jones, "Metallurgical Considerations in Linepipe Production",

American Welding Society Conference on Pipeline Welding and Inspection,

Houston, September 1982.

3. M. A. Civallero, C. Parrini and N. Pizzimenti, "Production of LargeDiameter High Strength Low Alloy Pipe in Itl ay". Microalloying 75

Conference, Union Carbide, Washington, D.C. 1975.

4. H. Nakasugi , H. Matsuda and H. Tamehiro, "Ultra-Low Carbon Bainitic

Steel for Linepipe", Conference on "Steels for Linepipe and Pipeline

Fittings", Metals Society, London, 1981.

5. F. Heisterkamp and K. Hulka, Niobium Products Company Ltd., internal

report, 1982.

P36-6
-T>6l£’ \ .

Chemical composition and mechanical properties specification for HY80 steel plates (from

MIL-S-16216G ships)

Chemical composition

Element

Per cent

Not less than Not more than

Carbon - 0. 18

Manganese 0. 10 0. 40

Phosphorus - 0. 025

Sulphur - 0. 025

Silicon 0. 15 0. 35

Nickel 2. 00 3. 25

Chromium 1. 00 1. 80

Molybdenum 0. 20 0. 60

Residual elements

Titanium - 0. 02

Vanadium - 0. 02

Copper - 0. 25

S + P < 0. 045%

Tensile properties

Thickness

mm (in.

)

range

0. 2 per cent proof stress,

0*0

Elongation on 50mm (2in.

)

gauge length, %

Reduction in area, %

Minimum Maximum Minimum Minimum

Less than 16 (0. 625) So loo 19 -

16 (0. 625) over

and

So - 95 20

55 longitudinal

50 transverse

Impact properties

i. Minimum energy absorption in Charpy V notch test Plate 11. 5mm (0. 5in. ) to 50mm (2. Oin. ) incl. 68J (50ft lb) at -84°C

No single specimen < 61J (45ft lb)

Plate over 50mm (2. Oin. ) 41J (30ft lb) at -84°C

No single specimen < 34J (25ft lb)

P36-7
Table Z. Carbon equivalent, weiaaDincy ana pre-neat calculations tor

HY 80 and similar steels (Ref. 1).

Preheat temperatures to avoid HAZ cold cracking (NCRE approach)

Carbon

equivalent, %

Thermal severity number (TSN)

4 8 12 16 20 24

0. 4

_ 40°C 70°C 85°C 95°C 100°C

0. 5 40°C 65°C 90°C 105°C 115°C 120°C

0. 6 65°C 90°C 110°C 125°C 135°C 140°C

0.7 90°C 110°C 130°C 140°C 145°C 150°C

0. 8 110°C 125°C 140°C 155°C 155°C 160°C

Comparison of hardenability of HY80, Q1 (N) and QT35 steels using carbon equivalent „ Mn Ni + Cu Cr + Mo + V CE% = C + — +

15 +

5

HY80, % Q1(N), % QT35, %

Typical ladle analysis for plate less than 44mm (1. 75in. ) thick 0. 71 0. 71 0. 64

Maximum specified ladle analysis for 44mm (1. 75in. ) thick plate 0. 95 0. 95 0. 75

Calculated minimum preheat temperatures to avoid HAZ cold cracking in plates having

typical analyses, assuming a restrained weld of thermal severity 18 and poor fit -up.

Carbon

equivalent,

%

NCRE

approach

,

°C

Welding Institute

approach

,

°C

UK MOD (Navy)

recommended preheating range, °C

QT35 0. 64 136 - 5 133 - 10 120-150

HY80 0. 71 144 - 5 156 - 10 120-150

Q1 (N) 0. 71 144 - 5 156 - 10 120-150

Table 3 . Typical Chemical Compositions of X-70 Pipe

Steel Type Gage C Hn Si_ fto Hi Cu Cr to V jn N B

Nb-V 18 0.09 1.50 0.15 0.16 0.10 0.12 0.04 0.08 0.016 0.005 - Mn-Mo-Nb 22 0.05 1.70 0.20 0.26 - - - 0.06 - 0.018 0.001 - Mn-Nb-V 22 0.08 1.90 0.27 0.24 0.20 0.24 - 0.05 0.06 . 0.007 - Nb-NI (14) 18 0.07 1.60 0.20 - 0.55 - - 0.11 - 0.019 0.004 - Nb-V-Cr 15 0.09 1.52 0.22 . - - - 0.35 0.04 0.08 . 0.008 - Nb-Mo-Cr 15 0.08 1.50 0.20 0.18 - - 0.30 0.05 - . 0.010 - Mo-Nb-Cu 15 0.06 1.45 0.20 0.25 - 0.30 - 0.025 - . 0.009 - Ti-Ni 15 0.08 1.58 0.30 - 0.26 - - - _ 0.070 0.004 - Ti-Cu-Nl (10) 20 0.09 1.61 0.30 - 0.12 0.21 - - - 0.063 - Low C-Nb 20 0.02 1.78 0.13 - - . - 0 . 10 - 0.015 0.004 - Low C-Nb-B (15) 20 0.02 1.89 0.13 - - - - 0.05 - 0.016 0.002 0.001 Low C-Nb-Mo-b (15) 22 0.02 2.01 0.16 0.30 - - - 0.05 - 0.018 0.002 0.001 C-Mn-Nb (0 4 T) 20 0.09 1.37 0.27 0.16 1.00 - - 0.02 - 0.020 0.005 - C-Mn-Nb 20 0.08 1.55 0.60 - - - - 0.06 - - 0.010 - {Controlled Cooled)

*0.25 for gages above 18-20 inn.

All Aluminum levels .02 - .05

All Sulfur levels .002 - .007

P36-8
Table 4" Typical Chemical Compositions of ULCB Steel

for Various Grades

U)

Steel Grade C Si Mn P S , Ni Mo Nb Ti B c# P *2J

P CM

UB-1 X - 65 0.02 0. 14 1.59 0.018 0.003 - - 0.04 0.017 0.001 0.29 0.11

UB-2 X-65 0.03 0. 16 1 .61 0.016 0.003 0. 1 7

- 0.05 0.016 0.001 0.31 0.12

UB-3 X- 70 0.03 0.14 1.91 0.018 0.003 - - 0.05 0.018 0.001 0.34 0.14

UB-4 X-70 0.01 0. 15 1.87 0.022 *3) 0.007

- - 0.04 0.020 0.001 0.32 0.11

UB-5 X-80 0.02 0.26 1.95 0.022 0.003 0.38 0.31 0.04 0.019 0.001 0.43 0.16

*1 ) Ceq

*2) P

CM

Cr*Mo+V Ni+Cu

5

+ 15 Mn+Cu+Cr Si V Mo Mi

C + + + + + — + 5B 20 30 10 15 60

*3) High-sulfur content

Table S Mechanical Properties of Pipe Body

Steel Grade Pipe

size

(mm)

Di rect

of

test

Tensile properties

Charpy V-notch

Impact properties

BDWTT

Y.S.

(MPa)

T.S.

(MPa)

El.

in 50. 8mm

(X)

Yield

ratio

(X)

Energy

at -20°C

(J)

501 Shear

FATT

(°C)

Shear fracture

area at -20°C (X)

80"* Shear

FATT ("0

UB-1 X-65 1420 00

x 16 WT

T rans- verse

500 583 42 85 369 < -80 100 -50

UB-2 X-65

1220 00

x 25 WT

T rans- verse 493 602 45 82 324 < -80 95 -30

UB-3 X-70

1420 0D

x 20 WT

T rans- verse

542 641 42 85 206 < -80 100 -40

UB-4 X-70

762 00

x 22 WT Transverse

551 622 41 89 159 < -80 98 -35

UB-5 X-80 1420 00

x 20 WT

T rans- verse

653 732 33 89 178 < -80 100 -50

MoVc : 7o Ws ; 5 4-?3 fhPov • SoWs; = 55o to/k

; Qolcs, =• C\<\ HP* ; loo fas M

Table G Mechanical Properties of Submerged-Arc Weld Joint

Transverse weld- Charpy V-notch

Pipe

size

tensi 1properties impact properties

Steel Grade

T.S.

(MPa)

Location

of

fracture

Test

position

Energy

at -20°C

(J)

50% Shear

FATT

(°C)

UB-1 X-65

1420 00 595Base

Weld metal 164 -40

x 16 WT metal HAZ 248 -50

UB-2 X-65

1220 0D

617

Base Weld metal 191 <-60

x 25 WT metal

HAZ 272 1-40

UB-3 X-70

1420 00

665

Base Weld metal 148 -40

x 20 WT metal HAZ 182 -35

UB-4 X-70

762 0D

633 Base Weld metal 183 -45

x 22 WT meta 1 HAZ 62*) -20

UB-5 X -80

1420 00

727 Base Weld metal 189 -60

x 20 WT meta 1 HAZ 130 -20

) Low energy due to high-sulfur content P36-9
Table 7. Chemical Composition and Mechanical Properties of

0.5% Ni , 0.1% Nb Steel

c Si Mn P S A1 Ni Nb

0.077 0.36 1.56 0.017 0.005 0.02 0.47 0.13

Y.S. UTS C

v (+20°C) 50% CVN-FATT 85% BDWTT-FATT

70.1 ksi 84.5 ksi 78J -85°C -55°C

56 inch diameter SAW pipe, 17.5 mm wall thickness

P36-10
X 70 H5LA LiNEPiPE STeELS

8*3* *ryy

0 07%c , , o^Nl

-+- OOT^V

OlO^Cr

2<* 03

Ult.mate tCMtae STSCM^TM ^Icsi

)

P36-11
X “Tou^UrxesS properWps »f Nc- Nib

Sutler

io^d«X ^'fe (l7*S *»*% u).t\).
Opportunities for Surface Modification Technology

in Conservation of Chromium

by

Peter G. Moore*

Although there are many other uses for chromium, three principal uses

are to improve the corrosion-, oxidation- or wear-resistance of structures

and materials. Because these properties are surface related, coatings can

be used in many cases to reduce the amount of chromium used and to allow the

surface properties of structures (as opposed to materials) to be tailored to

surface requirements without sacrificing bulk properties. A number of directed

energy team (DEB) processing techniques have been developed in recent years

which provide new opportunities for the conservation of chromium through reduced usage and through substitution for chromium by other elements. Today,

I will just describe in general terms two classes of processing techniques

(ion beam processing and laser beam processing) which produce novel corrosionand wear-resistant surface layers and then describe several cases in a bit

more detail.

With ion beam processing, the surface of the structure is bombarded by

ions with energies in the range of 10 to 200 keV. These ions are stopped within

1 pm of the surface of the solid material as a result of a series of collisions.

These collisions alter the microstructure of the surface layer and the implanted

ions alter the composition of the surface layer. This ion implantation requires

the production of an ion beam of the desired species. Ion mixing and ion plating

are variations where a thin layer of the desired material is first coated onto

the structure and then an argon or xenon ion beam is used to intermix the coating

and the substrate atoms.

* Peter G. Moore is with the Naval Research Laboratory

P37-1
With laser processing, the surface of the structure is rapidly heated by

a high power laser beam until the surface melts. After the beam is turned off,

the liquid rapidly resolidifies and is quenched as the heat is conducted away

to the bulk of the structure. Variations on this laser-surface melting process

include shock and transformation hardening (where no melting takes place) , surface alloying and cladding (where alloying elements are added) and particle inject Cohere wear resistant particles are incorporated into the melt) . There

are also many analogous processes using electron beams.

Directed energy beam (DEB) processed coatings in general have several advantages over traditional coatings: (1) a wide range of controlled coating thicknesses is possible, (2) coatings are metallurgically bonded to the subs tract,

(3) there are few thermodynamic constraints for laser processing and fewer still

for ion beam processing, and (4) structures and properties can be graded with

depth from the surface in order to optimize performance.

There are also several characteristics of DEB processing which restrict its

use. Almost all DEB processing is either done in vacuum or in a helium which

add to the cost of processing. Because all of the techniques are line-of-sight

processing, structures which can be processed by such techniques are restricted

to geometries which are relatively open and free of recesses. In addition, processing is limited to relatively small areas and structures, because of the vacuum

requirement for ion beam processing and because of warpage due to residual stresses

for laser beam processing.

Ion Implantation

As an example of the potential for solving materials problems through the

use of ion implantation, consider the study of the effect of implanting various

species on the resistance of M50 bearing steel to the pitting corrosion which

results in the salt water contaminated aircraft lubricants.

P37-2
Test specimens of M50 steel were implanted with 150 keV chromium ions to

obtain a peak chromium concentration of 24% at a depth of about 400 A below

the surface. Electrochemical polarization measurements and simulated field

service tests indicated an improved resistance to pitting corrosion and as a

result, an improved field performance.

Laser Processing

Three examples illustrate the potential of laser processing for solving

materials problems in the areas of corrosion, oxidation and wear. These are

the surface melting of 304 stainless steel to improve resistance to localized

corrosion, the surface alloying of chromium into steel substrate to produce a

stainless steel surface layer, and the production of wear resistant layers by

the particle injection method.

Recent work has attempted to improve the resistance of 304 stainless

steel to pitting corrosion by laser surface melting. The resolidified material

exhibited a very fine dendritic subgrain structure which indicated cooling rates

in excess of 10^K/sec. These materials were electrochemically characterized

by potentiodynamic polarization experiments in 0.1 M NaCl which indicated an

improved resistance to pitting compared to wrought 304 and comparable to that

of 316 stainless steel. A number of possible reasons exist for this improved

behavior as a result of the redistribution of both major and minor alloying

elements through homogonization during melting and segregation during and

after solidification. At this time, it is felt that the improvement is due

to the eliminiation of large precipitates which cause the local passive film

to be less protective.

Chromium steel surface alloys have been produced on various steel substrates using a variety of processing techniques. The electrochemical behavior

P37-3
of these alloys has verified that the surface layers passivate in the manner

of stainless steels. The broader significance of this and the previous results is that a stainless steel surface layer, suitable to the chemical environment can be rapidly and efficiently produced on the surface of a steel

structure

A third type of laser processing, particle injection, is performed in

the fashion of surface melting except that hard, wear-resistant particles

are injected into the melt pool. This results in a metal-matrix composite

surface layer which is metallurgically bonded to the substrate. Abrasive

wear tests of layers produced by the injection of TiC and WC into aluminum-,

titanium-, nickel- and iron-base alloys result in performance during abrasive

wear tests which are comparable to that of specialty wear coating materials.

An especially good example of the potential for this technique is the production of a TiC-aluminum matrix composite on an aluminum substrate; a wear resistant surface combined with a high thermal conductivity high specificstrength base metal which cannot normally be considered for applications where

wear is a problem.

Conclusion

Ion implantation and laser processing are both attractive techniques which,

when applied to appropriately, can be used to solve wear, corrosion and oxidation problems. The examples presented here have illustrated situations where

chromium can be conserved by using these coating techniques, and performance

comparable to or better than that of bulk alloys using chromium is obtained.

In these applications, chromium was used in a more or less traditional manner

and the role of the chromium in enhancing the performance has been the same

as its role in conventional alloys. By the nature of these techniques, there

are also many possibilities for enhancing performance of structures by proP37-4
ducing novel materials. It is this two-pronge potential of DEB processing

technique which will make them invaluable in the event of critical shortages

of key materials.

P37-5

SALT BATH TREATING AS AN ALTERNATIVE

FOR CHROMIUM PLATING

Nitrogen/Oxygen Synergism

The potential for replacing chromium plating by an oxidized nitrided

surface (Fig. 1) evolved from a completely unrelated event. Actually

the concept of following a liquid nitriding treatment by quenching in an oxidizing fused salt was originally devised to destroy any slight

traces of cyanide developed during nitriding and the substantial amount

of cyanate present in the salt as the active agent (Fig. 2). The

corrosion resistance of the resultant combination diffusion and conversion

coating is truly a synergistic effect.

The Aerated Liquid Nitriding Process

A more thorough understanding of the surface produced can be achieved by a brief consideration of the original nitriding process. The aerated liquid

nitriding process was introduced into the United States in the late 1950's

and early 1960's. The treatment requires immersion in a cyanide/cyanate

fused salt at 1060°F (570°C) for a 60 to 180 minute period. The salt is

aerated by a sparging ring immersed in the bottom of the fused mixture.

Make-up salt which consists of sodium and potassium cyanide is added and

reacts with the bubbled air to produce sodium and potassium cyanate

according to the following reactions:

KCN + 1/2 0

2 = KCNO (1)

NaCN + 1/2 0

2

= NaCNO (2)

The formation of the cyanate compounds in situ in conjunction with a 45 to

50% cyanide content and a 45 to 50% cyanate content is the basis for U.S.

Patent No. 3,022,204. During the nitriding action on carbon steels a 0.0004 in (.01 mm) to 0.0006 in (.015 mm) compound zone is formed in con- junction with a 0.018 in (0.45 mm) diffusion zone, (Fig. 3). The compound

zone is composed of single phase epsilon iron nitride (Fe^N) (Fig. 4)

responsible for its very unique wear resistance and low J coefficient of

friction properties. The nitrogen dissolved in the diffusion zone (Fig. 5)

is responsible for substantial increases in endurance values in both notched

and unnotched ferrous structural parts. In addition, the single phase

epsilon iron nitride surface exhibits excellent corrosion resistant properties

and in many cases the nitriding treatment is performed with this factor as a primary consideration.

P38-1
- 2 -

Sub-Critical Heat Treatment

Aerated liquid nitriding has the unique characteristic of being a subcritical single heat treatment capable of increasing corrosion resistance,

endurance values and wear resistance. The development of the epsilon iron

nitride surface and the nitrogen diffusion at temperatures substantially

below the critical temperatures of structural steels permits the elimination

of excess material and the subsequent required finishing operations to com- pensate for growth and distortion during processing. The process is thus

ideally suited for treating cylindrical shapes such as cylinder liners and

rocker arm shafts, where final finishing costs can be excessive. This

capability of retaining dimensional stability is also of vital importance

where corrosion protection is dependent upon thin films developed by heat

treatment thereby precluding final machining.

Process Changes Required Because of Environment

For many years this specialized heat treatment was documented, tested and

specified for crankshafts, gears, rocker arm shafts, connecting rods, valves,

valve guides and numerous other engine and chassis parts in all types of

industrial applications. Environmental pressures eventually restricted the use of cyanide compounds for metal processing in certain areas and it became

necessary to develop a relatively cyanide free liquid nitriding process. At

the same time the stated goal was to be able to produce a single phase epsilor

iron nitride compound zone (Fig. 6) with a nitrogen diffusion zone that could

produce corrosion resistance, endurance values and wear resistance equivalent

to those obtained in the existing process. The records accumulated during

years of testing of the original process made interchangeability a desired

asset.

New Liquid Nitriding Process

The aerated liquid nitriding process introduced in the 1970's met these

requirements in all respects. Unfortunately, under certain conditions the

fused salt also had the capability of generating small amounts of cyanide

during processing, according to:

KCNO + 2FeO = KCN + Fe^ (3)

and therefore could not be specified as a complete cyanide free salt.

Overcoming this objection involved post treatment in an oxidizing media

which began as a standard fused sodium and potassium nitrate/nitrite mixture

and then finally evolved into a sodium and potassium hydroxide, carbonate

and nitrate composition capable of total destruction of all traces of cyanide

and cyanate, (Fig. 7). The hardness patterns (Fig. 8) and (Fig. 9), wear

characteristics (Fig. 10) of the quenched nitrided parts exhibited minimal

change but the corrosion properties were substantially increased (Fig. 11).

Laboratory analyses of the quenched parts revealed that the unexpected

corrosion resistant results obtained were not a function of the jet black

surface (Fig. 12) but a more complex combination of nitrogen, oxygen and

iron.

P38-2
-3-

Auger profiles substantiated the fact that the diffusion of oxygen into

the epsilon iron nitrided compound zone was responsible for the corrosion

resistance. Studies eventually revealed that optimum dwell times in the

fused quenching salt would result in maximum corrosion resistance, although

this work is still in progress. The complete mechanism involved in con- verting a liquid nitrided surface into an effective corrosion barrier in

carbon and alloy steels is not completely understood. Test results how- ever continue to indicate that the majority of situations favor the

substitution of this treatment for chromium plating and this trend should

continue if the projected costs and availability of chromium metal eventually

become a reality.

Surface Compatibility with Bearings and Seals

Aerated liquid nitriding in conjunction with oxidizing salt quenching may be

a viable substitute for chromium plating but the process as performed lacks

compatibility in contact with many non-metal lie seals and bushings. With

this in mind a third and fourth step was added to the treatment when the

finished product is to be subjected to the above conditions. Basically the

objection to a standard nitrided and quenched part is the surface residue

or roughened finish resulting from nitriding, which can be a lap in contact

with softer nonmetal lie materials. Fortunately this difficulty can be easily overcome by polishing. Method of polishing is completely optional providing

surface removal is uniform and restricted to 0.00005 in (0.0013 mm) per

surface. Automated fixtured polishing with 320 grit paper has been used on bearing and contact surfaces of rotating or sliding shafts. Vibratory

finishing is completely acceptable where configuration prohibits fixturing.

Following the polishing operation the parts are immersed once again in the

quench bath for an oxygen diffusion period of twenty minutes.

The effect of this sequence of operations is shown in the bar graph of (Fig. 13).

Starting from the outside left bar, the finished carbon steel has very low

corrosion resistance and a standard commercial finish. After aerated liquid

nitriding, the corrosion resistance has increased dramatically, but directly

above, the surface roughness has also increased substantially. The third set

of bars illustrates how fine polishing reduces the corrosion resistance slightly

but effectively produces a surface finish suitable for bearing contact.

Finally, requenching in the oxidizing salt bath develops the ultimate combination

of surface finish and corrosion resistance.

The final process with oxygen diffused nitrided surfaces operating against

elastomeric and nonferrous seals or bearings is as follows:

1. Aerated liquid nitriding in a cyanide free fused salt -- 60 to 120

minutes at 1070°F (577°C).

2. Quench in controlled rate oxidizing salt at 750°F (400°C).

3. Water rinse.

P38-3
-4-

4. Abrasive polish to desired surface finish with maximum stock

removal of 0.00005 in (0.0013 mm).

5. Reimmersion in controlled oxidation bath, dwell time 20 minutes.

6. Water rinse and oil.

Corrosion and Wear Evaluation

Any system that is developed as a substitute for a proven existing product

must be subjected to extensive laboratory testing prior to acceptance. Also,

the testing has more credence if not restricted to the confines of the

developers research facilities.

Automotive spool shafts that are machined from plain carbon steel are subjected

to corrosive environmental effects when operating in their under hood locations.

A United States manufacturer of these shafts evaluated chromium plating,

electroless nickel and the oxidized nitrided coating as suitable protective

treatments (Fig. 14). The oxidized nitrided surface after the normal 88 hour

test duration reflected less than a 5% corrosion area which is significantly

less than the 70% figure shown for both chromium plating and electroless

nickel plating. Continuation of the test on the nitrided shafts only (to

produce failure) developed a 336 hour corrosion area of 5% to 10% which is

significantly less than the results obtained from the chromium and nickel

plated shafts during a 72 hour period. The results of these tests which

were performed in accordance with ASTM B117 are summarized in (Fig. 15).

One type of corrosion test performed in the research laboratory of a supplier

of automotive shock absorbers is shown in (Fig. 16). Two rods with different

surface treatments, chromium plating and oxidized liquid nitriding, were subjected to an accelerated life test of 300,000 cycles with a side load

applied, followed by 96 hours of ASTM B117 salt spray test. The loading

imparts a high pressure contact zone between the 0-ring seal and the piston

rod in a limited area. The figure illustrates that the chromium plated rod

has numerous corrosion sites on the surface while the oxidized nitrided rod

shows little effect from the corrosive atmosphere. Smaller rods, chromium

plated in the assembly, and liquid nitrided followed by quenching, foreground,

are shown in (Fig. 17).

In any discussion involving hard chromium plating, wear resistance properties

are usually as important as corrosion characteristics. A program was recently

initiated to measure the relative resistance of hard chromium plated with salt

quenched aerated liquid nitrided engine valves, (Fig. 18). The test load was applied through a rotating carbide cylinder making contact with a stationary

valve stem, oriented at right angles to the cylinder axis. Tests were performed

on both the intake valves, AISI 1541 and the exhaust valve 21-2N. After a one minute test period using prescribed loads, the degree of wear was measured as loss of weight in the valve components (Fig. 19).

P38-4
-5-

Summary

In summary, the oxygen diffused nitriding process has the capability of

producing a surface comparable or superior to chromium plating with respect

to wear resistance and corrosion resistance in numerous industrial applications. Unlike chromium plating, there is no requirement for the utilization

of strategic materials. Also, unlike electroplating processes, there is no water or air pollution problems involved when the oxidized liquid nitriding

process is practiced. The materials required to produce both of the fused

salts involved in the process are readily available coironon chemicals and none of the compounds present in the effluent rinse waters are presently

categorized on a restricted discharge basis. Unit costs of any process

are usually suspect after a very short period of time because of our rapidly

changing economic conditions but on a comparative basis the total costs

involved in chromium plating and the oxygen diffused nitriding process are quite similar and are likely to be that way even with future changes in

the Cost of Living Index.

P38-5
Fig. 1. Ball studs, aerated liquid nitrided and

salt bath quenched.

Decomposition of Cyanide and Cyanate in

Quench Salt

CN- in mg/kg

i \ —\

1

\

Cyanide

Analyzed after 3 minutes

o 5 minutes

10 minutes - 30 minutes

L

"

:

CNO' in mg/kg

Cyanate

Analyzed after k 3 minutes

5 minutes

|_3 ! 10 minutes - 30 minutes

200 250 300 350 200 250 300 350 Temperature of Quench Salt in °C

Fig. 2. Chemical destruction of cyanide and cyanate

residues, after salt bath quenching, related

to time and temperature.
COMPOUND

ZONE

DIFFUSION

ZONE

CYANIDE-FREE LOW TEMP. LIQUID NITRIDED

90 MINUTES AT 1060F (570C)

Fig. 3. Compound zone and diffusion zone developed by

low temperature liquid nitriding. Note the

needle like structure of Fe^N developed in

the diffusion zone by aging.

Fig. 4. The epsilon iron nitride phase in aerated liquid

nitriding is developed, as shown, slightly to the

right of the intersection of the 8% nitrogen line

with the 1060°F (570°C) temperature line. The

stoichiometric ratio of nitrogen to iron is very

close to one to three.

P38-7
Fig. 5. Nitrogen distribution in the diffusion zone of

aerated liquid nitrided carbon steel. It is

this nitrogen that is responsible for increases

in endurance values from 20% to 100%.

Fig. 6. X-ray diffraction pattern of aerated liquid

nitrided carbon steel. The three typical

peaks of epsilon iron nitride are present

while the absence of other compound peaks

indicate that the zone is single phase.

P38-8
CN~

in

mg I

kg

Decomposition of cyanide and

cyanate in AK-Bath

Fig. 7. Chemical destruction of cyanide and cyanate in

an oxidizing media.

Fig. 8. Hardness pattern resulting from aerated liquid

nitriding of plain carbon steel. The sharp

drop below the surface measures the depth of the

compound zone.

P38-9
Fig. 9. Hardness pattern developed by aerated liquid

nitriding of alloy steel. The higher surface

hardness and the measurable retention through

the diffusion zone results from alloy influence.

0 12 3 4 RUNNING TIME IN HOURS

Fig. 10. Wear test comparison of untreated, carburized

and aerated liquid nitrided specimens.

P38-10
I

!

)

'

Fig. 11. Salt water corrosion tests on various surface

treatments. The SBN-SBQ notation refers to

salt bath nitriding followed by salt bath

quenching. The other treatments are standard.

Fig. 12. The top gas spring has an aerated liquid nitrided

shaft compared to the lower spring with a chromium

plated shaft.

P38-11
Fig. 13. The effect of sequential surface

treatments on surface finish and

corrosion resistance. The acronyms

for aerated salt bath nitriding,

SBN, and oxidizing salt bath quenching,

SBQ, are used.

Corrosion Resistance Evaluation

Spool Shafts

Test Conditions: ASTM B-117

QPQ Hard Chrome Plate

After 336 hours After 72 hours

14. Automotive spool shafts subjected to salt spray

testing according to ASTM B-117. The nitrided/

quenched surface is indicated by QPQ.

P38-12
Fig. 15. Resistance of various surface treated spool

shafts exposed to corrosion testing according

to ASTM B117. QPQ refers to the quenching and

polishing of aerated liquid nitrided specimens.

Corrosion Resistance

Piston Rods

Tsst Conditions: 300,000 cycles endurance tests followed by 96h salt spray, ASTM B-117 Hard Chrome

Plate QPQ

Failed; more than ten Passed; no evidence

corrosion sites per sq. In. of corrosion

Fig. 16. The aerated liquid nitrided piston rod after salt

bath quenching, polishing, and requenching exhibits

no evidence of corrosion after endurance testing

followed by 96 hours of salt spray.

P38-13
Fig. 17. Shock absorber assembly and rod. Combinations

of endurance properties and corrosion resistance

are required.

Hard Chroma

Plated QPQ Treated

ENGINE VALVES

Fig. 18. Surface treated engine valves.

P38-14
Wear Resistance Evaluation

Surface Treatments

Test Conditions: Engine valve stems, cross cylinder wear test, 1 minute duration Log

Weight Loss

(grams) 2000 -i

I SBN - SBCL

I Chrome Plal

Mil 1 l Zero Load 25 lb. Load Zero Load 25 lb. Load

Intake Valve Stem Exhaust Valve Stem

Fig. 19. Comparison of the wear properties of hard

chromium plated engine valve stems with

stems treated by aerated liquid nitriding

followed by salt quenching.

P38-15

CLAD METALS - MATERIAL CONSERVATION THROUGH

DESIGN FOR CORROSION CONTROL AND HIGH PERFORMANCE

James T. Skelly

Metallurgical Materials Division

Texas Instruments Incorporated

Attleboro, MA 02703

Abstract

Growing concern over the continued availability of critical

materials at a reasonable cost has spurred the development of new

materials, product redesign, and improved conservation and

recycling methods. The pursuit of new materials that reduce or

even eliminate reliance upon critical materials can offer

tremendous payoffs. Numerous research facilities are

concentrating on new alloy development, surface coatings,

surface modification, ceramics, and composites, all of which hold

promise in reducing dependence upon unstable foreign sources for

critical materials. Clad metals provide another approach.

By precisely engineering a clad metal system, unique

functional benefits can be achieved that are otherwise

unavailable or available only with heavy reliance on critical

materials. This paper will focus on the design, development,

application, and performance of precision clad metals. Specific

unique properties and corrosion control benefits will be

discussed as will feasible alloy combinations.

Cladding Technology

The modern process for production of clad metal strip is

focused upon continuous cold roll bonding of two or more strips.

Prior to bonding, the individual strips of metal are rolled to

precise gauges and extensively cleaned to provide contaminantP39-1
free surfaces. When passed through the specially-designed high

pressure rolling mill, a composite material is formed, deriving

its bond integrity from the merging of the atomic lattices of the

two metals into a common structure. Subsequent thermal treatment

is performed to promote diffusion, improve bond strength, and

provide stress relief for further cold work operations. Finishing operations, such as rolling to intermediate and final gauge,

annealing, buffing or polishing, and slitting ensure matching

clad metal properties to specific and exact product requirements.

These precision clad metals are manufactured by a highly

refined process which is based upon the technology of solid state

welding. No intermediate brazing alloys or adhesives are used to

achieve a permanent bond.

It is important to distinguish clad metals from plated

metals. Most importantly, the layers in the clad metal system

are wrought metal as opposed to thin and porous plating layers.

Also, the thickness of the individual clad metal layers can be

precisely controlled over the entire composite metal thickness as

opposed to the well-known practical limitations to plating

thickness. Figure 1 compares the qualities of cladding versus

plating.

Generally, most metals and alloys can be clad. Figure 2

illustrates that many ductile metals and alloys available in

t wrought form are in production today. Numerous other

combinations are technically feasible but require development

before commercialization.

The Clad Metal Approach

At first glance, direct substitution of a clad metal system

P39-2
to conserve critical materials can be envisioned as a surface

layer of an alloy containing critical materials clad to a widely

available, low cost base metal. In fact, origins of modern clad

metal manufacture addressed conservation of precious metals such

as gold or silver.

Direct substitution of clad metals for critical materials,

however, can only offer a partial solution. The usage of the

critical material is only reduced and not eliminated in the end

product. The process to clad the critical material will offset

much of the apparent economic advantage of the critical material

content reduction. We should not be looking for substitutes but

should be looking for alternate approaches.

Optimal application of clad metals for conservation of

critical materials is derived from obtaining a longer useful life

via corrosion control or improved performance of the end product

without the heavy reliance upon critical materials.

Metals Design for Corrosion Control

Controlling the effects of corrosion with the use of clad

metals necessitates utilization of the galvanic series, which is

based upon the electrochemical potential of a metal and a reference cathode. Figure 3 illustrates a typical galvanic

series in seawater where a less noble metal would corrode at

accelerated rates when coupled electrically to a more noble

metal. The application of the galvanic series in design of clad

metal systems can result in enhanced corrosion control

performance. The corrosion control mechanisms and example

applications will be discussed below.

1 . By cladding a thin layer of highly corrosion-resistant

P39-3
metal over a less corrosion-resistant and lower cost metal,

it is possible to achieve enhanced and reliable corrosion

resistance. The less noble core layer provides the mechanical strength at a low cost. The system differs from plating

in the thickness and nonporosity of the noble layer.

An example of this approach, called noble metal

cladding, is stainless steel clad aluminum for truck

bumpers. The highly corrosion-resistant type 301 stainless

steel is 10 times the thickness of chromium plating over a steel or aluminum substrate. As shown in Figure 4,

laboratory and field testing has confirmed that defects in

the chromium plating layer or damage during installation or

service which results in a small anode to cathode ratio, can

lead to shorter life for a conventional plated steel or

aluminum bumper. The stainless steel clad aluminum bumper

has demonstrated beyond a reasonable doubt that it can meet

the m i 1 1 i on-m i 1 e- 1 i f e demands of the heavy duty truck

bumper

2. The combination of a more noble metal clad to a less noble metal, where the purpose of the less noble metal

is to protect the more noble metal by corroding

preferentially, is an example of a sacrificial corrosion

system. This principle has been widely applied in the

automotive industry where stainless steel clad aluminum trim

has replaced solid stainless steel trim. As illustrated in

Figure 5, the attachment of stainless steel directly to the

automotive body will promote accelerated corrosion of the

body steel. Conversely, as shown in Figure 6 the aluminum

P39-4
cladding protects the automotive body from corrosion, by

corroding preferentially to both the steel and the stainless

steel, thereby extending the life of the vehicle.

Solid aluminum trim could provide similar corrosion

protection of the body steel, but aluminum itself is far

less corrosion resistant than stainless steel, and would

rapidly oxidize, detracting from the vehicle’s appearance.

3. By cladding a less noble layer over a more noble layer,

a corrosion barrier is formed where the more noble inner

layer begins to corrode only after the outer layer has been

extensively corroded. By designing sufficient mechanical

strength for the application into the inner layer, a material system has been created that greatly exceeds the

corrosion resistance of the individual components.

A clad system of carbon steel clad to both sides of a

core of stainless steel for automotive brake tubing

demonstrates this corrosion barrier mechanism. Figure 7

compares the corrosion behavior of carbon steel tubing and

carbon steel clad stainless steel tubing. While the conventional terne coated carbon steel tubing deteriorates rapidly

due to localized corrosion, the corrosion in the clad tubing

is arrested at the stainless steel layer due to the galvanic

protection provided by the carbon steel.

4. As an insert between dissimilar metals, a clad transition material eliminates the opportunity for destructive

corrosion at the joint. It also simplifies attachment by

welding. The joining of aluminum to steel, difficult by

welding and doomed to crevice corrosion failure, can be

P39-5
accomplished by placing an insert of aluminum clad steel.

The aluminum component is welded to the aluminum layer and

the steel is welded to the steel layer. The presence of the

metallurgical bond eliminates the opportunity for crevice

corrosion.

5. For a material system to withstand differing corrosive

environments on each side of a strip requires a monometal or

( alloy to resist the corrosive effects of both environments.

An alternate approach is to develop a clad metal system with

each layer designed to resist the corrosive effects of the

individual environment.

An example of a complex multilayer clad system has

been developed for anode caps for battery button cells. The

inner copper layer is compatible with the internal chemistry

of the battery cell while the outer nickel layer resists the

corrosive effects of the external environment. The middle

stainless steel layer provides the formability and

0 mechanical strength.

As illustrated in Figure 8, another example is a three-layer system of titanium clad copper clad nickel for a bipolar electrode in fuel cells. The nickel is required for

its resistance to hydrogren permeability on the cathode

side. The highly corrosive environment necessitates titanium

on the anode side. The copper core is required for its

thermal and electrical conductivity to carry electrical

current and dissipate heat.

Through appropriate design, the five clad metal approaches to

corrosion control described above can greatly extend the useful

P39-6
life of end products. This extended life itself can have

dramatic reduction in the demands for critical materials.

Clad Metal Design for High Performance

In a similar fashion to corrosion control with clad metals,

the combination of materials into a clad metal system can often

result in a unique combination of physical properties or

properties that might otherwise be achieved only through the use

of expensive or critical materials. Properties, such as

strength, ductility, electrical and thermal conductivity, thermal

expansion, magnetism, and surface appearance can be tailored to a

specific requirements in a clad metal system with exacting

precision

A key tool in the development of clad metal sytems is the

rule of mixtures, which is a series of calculations used to

determine the new composite’s properties, based upon the known

properties of the individual metals. The rule of mixtures

equation can provide a good approximation of the composite’s

properties in order to narrow the field of potential clad metal

systems during the design phase, without actually fabricating the

composite

The general equation which can be applied to calculate

properties such as density, lateral thermal conductivity,

electrical conductivity, and thermal expansion through the thickness of the composite is shown below:

D =Ai x D i + A 2 x D 2 + .... An x Dn

Too

Where A-], A2> etc., are the thicknesses, in percent, of the

P39-7
components of the clad metal.

D -| , D 2 > etc., are properties of the same components.

Properties such as thermal conductivity normal to the

surface, thermal expansion along the length, and modulus of

elasticity can be calculated with different equations. For

fatigue and mechanical strength properties, testing of actual

sample specimens or prototypes is recommended.

1. An example of functional benefits offered by a precision

clad metal is high thermal and electrical conductivity

copper clad to both sides of a low thermal expansion rate

64$ iron-36$ nickel alloy known as Invar. As illustrated in

Figure 9> it is being successfully applied as a printed

wiring board core material for direct surface mounting of

ceramic chip carriers.

In order to achieve more dense electronic packaging,

ceramic chip carriers were developed to replace the existing

dual inline packages (DIP), thereby reducing component area

by factors of 2-5:1. Because the solder joint is both the

mechanical and electrical attachment medium for these

components, an exact match is required in thermal

coefficient of expansion (TCE) between the ceramic chip

carrier and the printed wiring board. This is especially

critical when wide operating temperatures are encountered

such as in military, automotive, and telecommunication

applications. Also, the greater packaging density achievable

with chip carriers can generate more heat per unit area,

which needs to be dissipated to prevent high temperature

failure of the integrated circuit. Additionally, strength,

P39-8
stiffness, and flatness are key criteria in selection of a printed wiring board material system.

Unlike other material systems, copper clad Invar offers

a combination of properties that allow the electronics

packaging engineer to achieve all the required performance

levels. Copper clad Invar is now being designed into applications which had previously utilized cobalt-containing

Kovar due to its low thermal expansion rate. With copper

clad Invar, it is possible to achieve a thermal expansion

rate comparable to Kovar as well as superior thermal and

electrical conductivity properties. Figure 10 compares the

thermal properties for copper clad Invar to other printed

wiring board materials.

2. Another example of the synergistic attributes of clad

metals is the application of stainless steel clad aluminum

for aircraft firewalls. Titanium had been considered due to

its excellent strength-to-weight characteristics and its

high temperature performance. However, titanium's cost and

availability make it an undesirable choice.

Stainless steel clad aluminum provides strength-toweight properties similar to that of titanium and has passed

the stringent FAA flame test (15 minutes at 2000°F). Best

of all, this material is readily available at a fraction of

the cost of titanium.

3* Another high performance clad metal system is copper

clad to both sides of a ferritic stainless steel core for

integrated circuit lead frames. Copper clad stainless steel

lead frame material has the ductility and strength

P39-9
equivalent to or better than the traditional lead frame

material, Alloy 42 (nickel-iron alloy), yet provides six

times the thermal conductivity of Alloy 42.

With integrated circuit devices becoming increasingly

complex, they need improved thermal dissipation to avoid

high temperature failures. The copper layer provides the

excellent lateral thermal conductivity while the stainless

steel layer provides the strength and ductility. Copper

alloys, which are also being used in high power IC devices,

have marginal mechanical properties for many lead frame

applications

4. Heat exchangers for industrial and transportation applications have traditionally been made by inserting copper

shims between the structural layers and furnace brazing the

entire assembly. The approach can be highly labor intensive

in compactly designed units and can also lead to defective

braze joints due to improper location of the copper shims.

As illustrated in Figure 11, the clad metal approach

utilizes a composite of copper clad to both sides of a core

of carbon steel or stainless steel. The copper acts as the

brazing alloy and thus eliminates the cost of manufacturing

and handling the copper shims. The clad metal system also

ensures a sufficient presence of copper for requisite braze

joint strength.

Application of Clad Metals for Critical Material Conservation

New alloy development, which holds promise in reducing

dependence upon materials such as chromium, cobalt, manganese,

and titanium, requires up to ten years and often increases

P39-10
material cost or reduces performance.

Innovative design and application of clad metal systems can

sharply reduce the demand for critical material without

sacrificing performance and cost. Additionally, clad metal

development can take as little as one year. In most cases,

readily available materials can be combined into clad metal

systems that yield unique functional properties and corrosion

control capabilties.

When the U. S. Mint was searching for alternate material

systems to silver bearing coinage, cupronickel clad copper clad

cupronickel was developed in approximately one year. Furthermore,

the clad coins met strict Mint requirements for corrosion

resistance, formability, appearance, abrasion resistance, Rho

density, reclaimable scrap, and availabilty.

Innovative design and production of clad metals have

resulted in widespread use in numerous industrial and consumer

applications. The capability to engineer a clad metal to operate

in hostile, corrosive environments and meet exacting performance

requirements can be readily applied to the conservation of

critical materials. Our engineering community, designing

tomorrow’s products, must include potential clad metal systems

early in the design and evaluation phase of these new products.

When designed in from the beginning, the full payback in terms of

material conservation, optimal performance, and inherent

corrosion control can be achieved through clad metal technology.

Many engineers, unaware of the large volume availability and

advanced technology of clad metals, do not consider clad metal

systems but specify alloys that consume large amounts of critical

P39-11
materials. Often design engineers make trade-offs among

engineering properties by accepting some properties that are

marginal to retain other more critical properties, or by overdesigning the entire system often using critical materials. This

overuse or misuse of materials can led to marginal performance of

the end product, shortened life in the operating environment as

well as increased cost.

An example of this is the use of Kovar (17% Co/29% Ni/

balance Fe) in numerous electronic components utilizing glass-tometal seals due to its low TCE and glass-matching capability up

to 450°C. Kovar, however, is extremely costly, consumes precious

cobalt, and is a poor thermal conductor. Copper clad Invar

offers a similarly low TCE through 225°C, improved thermal

conductivity, a solderable surface, and eliminates the

consumption of cobalt. Process modifications to allow copper

clad Invar usage might necessitate a shift to a solder seal

approach to accommodate the transition of Invar to a higher

expansion rate at 225°C. However, it is these kinds of process

changes that can reduce demand for critical materials with the

added benefit of reducing cost.

Summary

Material conservation begins with the product design and

evaluation stage. The design engineer must comprehend not only

the initial product design but also its life cycle-performance

requirements. The design engineer must seek innovative designs

and processes that reduce dependence upon critical materials. The

competitive marketplace demands such innovation; the motivation

is lower cost as well as improved performance.

P39-12
COMPARATIVE QUALITIES OF

CLAD VERSUS ELECTROPLATED COATINGS

Clad Laminate

Advantages

Clad layers have low porosity.

Composition is precisely

controlled

Alloys and active metals

(Al, Ti, etc.) can be clad.

Thickness of cladding is

continuously variable and

precisely controlled.

Properties do not vary with

cladding thickness.

Complex, multilayer systems

can be produced.

Limitations

Fabricated parts cannot be

clad

Cut edges may require special

treatment

Brittle metals cannot be clad.

Width of stock is limited.

Electroplate

Advantages

Fabricated parts can be

plated

Brittle materials can be

plated

Finish brightness can be

controlled

Selective areas can be

plated.

Limitations

Deposited metal may be

porous

Codeposition (with S, H 2 , etc-'' is common.

Active metals (Al, Ti,

etc.) cannot be deposited.

Substrates with stable

oxides cannot be plated.

Thickness is limited by

plating time.

Properties can vary with

plating thickness.

Thickness and uniformity

can vary.

Figure 1

P39-13

GALVANIC SERIES IN SEAWATER

NOBLE END

Platinum

Gold

Silver

Titanium

316 Stainless Steel

304 Stainless Steel

410 Stainless Steel

Nickel (passive)

Cupronickel

Bronzes

Nickel (active)

Tin

Lead

316 Stainless Steel

304 Stainless Steel

Lead-Tin Solder

410 Stainless Steel

Cast Iron

Steel

Cadmium

Aluminum

Zinc

Magnesium

( passive

)

( passive)

(passive)

(active)

(active)

( active)

ACTIVE END

Figure 3

P39-15

,v%-

P39-17
P39-18
CYCLIC IMMERSION (DIP AND DRY) CORROSION TEST

OF BRAZED AND TERNECOATED BRAKE LINE TUBING

IN SIMULATED ROAD SALT ENVIRONMENT

INITIAL

20,000

CYCLES

40,000

CYCLES

Figure 7

P39-19
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P39-22
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References

1) Hodesblatt, S. , Alternatives to Strategic Materials, Design

Engineering, January, 1982, pp 44-49

2) Dance, F.J., Clad Metal Circuit Board Substrates for Direct

Mounting of Ceramic Chip Carriers, Electronic Packaging and

Production, January, 1982

3) Dance, F.J., Mounting Leadless Ceramic Chip Carriers

Directly to Printed Wiring Boards - A Technology Review and

Update, Nepcon West, February 24, 1982

4) Bowers, E. W., the U.S. Engaged in Another Kind of War - Strategic Minerals, Iron Age, April 14, 1982, pp 38-41

5) Hurlich, A., Strategic Materials - Technology Trends,

Mechanical Engineering, July, 1982, pp 44-53

6) Baboian, R., Designing Clad Metals for Corrosion Control,

SAE Trans., 8j_, pp 1763, 1973

7) Baboian, R., Clad Metals in Automotive Trim Applications,

Paper 710276 SAE Automotive Engineering Congress, January,

1971

8) Baboian, R., Corrosion Resistant, High Strength Clad Metal

System for Hydraulic Brake Line Tubing, SAE Trans., 8

, pg 1117, 1973

9) Baboian, R., and Haynes, G., Joining Dissimilar Metals with

Transition Materials, paper 760714, SAE Automotive Engineering Congress, January, 1976

10) Baboian, R. and Haynes, G., Corrosion Barrier Materials for

Communication Industry, Materials Performance, J_6 , pg 30,

1977

11) Baboian, R., Controlling Galvanic Corrosion, Machine Design,

October 1 1 , 1979

12) Delagi, R., Designing with Clad Metals, Machine Design,

November 20, 1980

13) Regan, R., Metals Industries Waiting for Reagan's Minerals

Policy, Iron Age, December 16, 1981, pp 31-33

14) Cassidy, R., Strategics: The New Gold?, Sky, August, 1981,

pp 10—16

15) Baboian, R., Clad Metals Respond to a Changing Automotive

Environment, Body Engineering, pp 69-74, Spring, 1981

16) Verink, E.D., Our Next National Crisis: Materials,

Mechanical Engineering, pg 42, August, 1980

P39-24
17) Baboian, R., Conservation of Critical Materials with Clad

Metal Systems, Electrochemical Society Transactions

18) Gomez, P. J., A Systems Approach to the Economic Use of

Natural Resources, Wire Association Convention, October,

1966

19) Quimby, W. G. , Production Methods, Properties and Applications for Clad Metals, American Society for Metals Meeting,

March 21 , 1979

P39-25
Autobiography

James T. Skelly is a graduate of the University of Rhode

Island where he received his B.S. in Industrial Engineering. He

also has an MBA from Northeastern University. He has been

employed by Texas Instruments in finance, production control in

the Control Products Division, and automotive program manager for

the Metallurgical Materials Division. Jim is currently product

manager in the Metallurgical Materials Division.

P39-26
ABSTRACT

It has long been known that carbon steel exposed to geothermal brine

is aggressively attacked and large corrosion allowances must be made in the

design of vessels and piping used in such environments. The economics of

geothermal power presents a real need for the use of functional low-cost

materials as liners for exposed carbon steel surfaces.

Polymer concrete (PC) has been identified as a promising liner

material. PC is defined as a concrete in which the aggregate is bound in a dense matrix with a polymer binder. Several high-temperature PC systems

have been formulated and tested in the laboratory in brine, flashing brine,

and steam at temperatures up to 260°C. Results are also available on field

exposures lasting up to 36 months from at least one of six geothermal test

sites. Good durability is indicated in all tests.

A study has indicated that the use of PC liners as a replacement for

the corrosion allowance in carbon steel components for a 50-MWe geothermal

plant will reduce the cost of electrical generated power by ~6.2 mills per

kilowatt hour.

P40-1
INTRODUCTION

The demand for the development of geothermal energy increased as the

transitory era of cheap imported fuel came to an end. The constant

increase in global energy consumption has made even greater demands for the

production of a cheaper energy source.

One such source being investigated is the extraction of heat energy

from the earth, i.e., geothermal energy. Worldwide commercial utilization

of geothermal energy was reported to be over 7000 megawatts in 1979. ^ A

1980 survey performed by EPRI indicated that current geothermal electricity

generation in North America was 663 megawatts and projected a growth to

about 10,000 megawatts by the year 2000. In addition, potential direct

utilization of geothermal energy by a wide range of industries was also

indicated .

The development of geothermal energy has encountered some serious

technical problems in the handling of hot brine and steam. Hot brine and

other aerated geothermal fluids are highly corrosive and they chemically

attack most conventional construction materials. Corrosion and scale

encrustations have been encountered in all geothermal plants, and to

various degrees they adversely affect plant lifetimes and power output. To

date, carbon steel has been the primary material of construction but the use of expensive materials such as chrome-moly steel and titanium base

alloys may be required for long term operation. Corrosion studies of

several different steels exposed to high-temperature geothermal brines were made by the Lawrence Livermore Laboratory at the Salton Sea Geothermal

Field. ^ Six-month exposure studies indicated corrosion rates of 1.65

mm/year for mild steel (AISI-1009) while chrome-moly steel (ASTM-A387-9

)

had a corrosion rate of 0.13 mm/year. Chrome-moly steels, however, cost

about four to six times as much as the carbon steels, thereby substantially

increasing plant construction costs.

P40-2
The availability of low-cost materials with suitable properties will

enhance the development of a stable geothermal industry. Brookhaven

National Laboratory (BNL) under contract tc the Department of Energy (DOE)

has been developing non metallic construction materials suitable for use with geothermal fluids. Calculations indicate that if successfully

implemented, the use of materials such as polymer concrete (PC) in geothermal electrical generating processes could result in a 10 to 20% cost reduction. On the basis of projections of USA direct utilization applications

by the year 2000, energy cost savings equivalent to an annual saving of

$ 50x10 ^ to $ 300x10 ^ have also been estimated.

Development of non metallic construction materials for use in geothermal systems was initiated in the early 70s. Since that time, several

high- temperature PC systems have been formulated. Laboratory and field

tests have been performed in brine, flashing brine, and steam at temperatures up to 250°C. Laboratory data for exposure times of up to 2 years are available. Field test data from six geothermal sites with exposures of up

to 3 years are also available. Durability of the PC has been good in

almost all cases. The results to date have indicated the potential for the

successful use of a specially formulated PC as a lining material for

process piping and vessels in geothermal power systems. PC glass filament

wound pipe, with low thermal conductivity, can be made for process heat and

district heating applications.^

Economic studies performed concurrently with the research program have

identified several cost-effective uses for PC in geothermal processes. One

study has indicated that the use of PC liners as a replacement for the

corrosion allowance in carbon steel components for a 50-MWe geothermal

plant will reduce the cost of power by ~6.2 mills per kilowatt hour.

Greater savings are indicated if PC is substituted for stainless steel,

titanium, or Hasteloy in acid-handling systems, condensate-piping systems,

reinjection lines, and steel separators. Uses in cooling towers, district

heating systems, and the protection of concrete surfaces also appear to be

cost effective.

P40-3
PRODUCTION METHODS

Polymer concrete is defined as a concrete in which the aggregate is

bound in a dense matrix with a polymer binder. The techniques used for

mixing and placement are similar to those used for port land cement con- crete, and after curing a high-strength durable material is produced. The

most important process variables are monomer and aggregate composition and

the aggregate particle-size distribution. Specimens can be produced with

compressive strengths up to 207 MPa. Full strength is attained immediately

after the polymerization reaction is completed. Polymerization can be

accomplished with initiators and promoters at ambient temperatures or with

initiators in conjunction with heat. Combinations of both methods are also

used to ensure complete polymerization of the monomers used.

A. Monomer Formulations

Several monomer systems that can be used in high-temperature PC formulations have been developed. Some of these will be briefly reviewed.

Long-term test data are available for the following systems:

1. 60 wt% styrene, 40 wt% trimethylolpropane-trimethacrylate (TMPTMA)

2. 50 wt % styrene, 33 wt% acrylonitrile, 17 wt % TMPTMA

3. 55 wt % styrene, 36 wt % acrylonitrile, 9 wt % TMPTMA

These systems can be polymerized using chemical initiators (benzoyl

peroxide, di-tert-butyl peroxide) and heat or by chemical initiators

(benzoyl peroxide) and promoters (dimethyl aniline). The styrene-TMPTMA

mixture is suitable for the manufacture of PC which is stable to 150°C,

while the last two systems can be used at temperatures up to 250°CB. Aggregate Selection

The durability of PC to geothermal fluids is highly dependent upon the

composition of the aggregate. Materials such as quartz, silica, flyash,

and portland cement have been investigated. All of the aggregates have

P40-4
been used successfully in materials at temperatures <210°C. Above this

temperature, only PC materials containing a mixture of silica sand and

Portland cement have been durable when subjected to brine and steam.

Experimental work at BNL indicated that composites formed with a mono- mer system containing styrene(S), acrylonitrile (ACN), and TMPTMA, in con- junction with an aggregate system containing silica sand and portland

cement (10 to 40% by weight of aggregate mixture), were stable in 25% brine

solutions and steam at temperatures up to 250°C* Bonding of the monomers

to the silica sand was enhanced by the use of a silane coupling agent. It was also observed that Type III portland cement gave better results than

Type I or IV. This led to the belief that a bond was formed with the vinyl

monomers and the tricalciura silicate (highest in Type III cement) which

resulted in increased thermal and chemical stability >-* Later experimental work at BNL verified that a bond does exist between vinyl-type

monomers and the cement phase thus cement-containing PC would be expected

to have higher thermal stability and long-term durability.

The compressive strengths of PC specimens tested at the Geysers,

California (Figure 1), and duplicated in laboratory autoclave tests show

that they decay for the first 30 to 60 days and then remain constant. Test

data have been collected for periods up to 2 years, and some of the data

are given in Table 1. Additional data collected from specimens exposed to

25% brine solutions at 177°C in laboratory autoclaves are given in Table

2. The data indicate that the strengths of the samples are essentially

constant after the first 60 days of exposure, with samples containing

acrylonitrile having higher compressive strengths.

C. Manufacturing

PC-lined pipes were manufactured at BNL and installed at the U.S.

Bureau of Mines Corrosion Facility at Niland, California, in the Salton Sea

geothermal field. This is considered to be the most aggressive corrosive

fluid known to date. It contains 280,000 ppm dissolved solids in the brine

at wellhead temperatures of 240° to 260°C. The lined pipe was installed in

P40-5
the inlet line of the corrosion facility and was in continuous service for

< months (Figure 2). The pipe was 1.2 m long and had an inside diameter of

7.6 cm.

The liner was made of PC containing 55 wt% S, 36 wt% ACN, 9 wt%

11PTMA. The aggregate system was 70 wt% graded silica sand, 30 wt% Type

fll portland cement. The inside of the steel pipe was sandblasted to remove any scale and/or rust. The inner core was made of rubber (7.6 cm diam) and installed in the center of the steel pipe to leave an annulus of

1.25 cm. Benzoyl peroxide 98 (1%) and di-tert-butyl peroxide (0.5%)

Initiators were dissolved in the monomer mixture. The aggregate system

( -90%) was loaded in a Day rotating blade mixer and the monomer mixture

( 10%) added. When the composite was thoroughly mixed the PC mortar was

p

1 aced in the annulus while the pipe was mounted on a table vibrator.

Longer or larger-diameter pipes would have additional vibrators on the pipe

wall. When the annulus was filled, the PC was allowed to cure in situ . Alter the initial cure, the inner rubber hose was removed and the lined

pipe placed in an oven at 150°to 190°C for a final cure.

Numerous steel pipes have been lined with PC and tested at various

geothermal sites. Pipe diameters have ranged from 7.6 to 30.5 cm and

lengths from 1.5 to 2.5 m. At East Mesa, California, PC-lined pipes

(Figure 3) were in service for up to 3 years with no apparent decay.

Polymer Concrete Research, Inc. under contract to BNL has developed a centrifugal casting process that can be used for the commercial manufacture

of PC-lined steel pipe.^ A steel shell is placed in the centrifugal

casting machine and using a lance the inside of the shell is cleaned by

sandblasting (Figure 4). Preweighed PC batches are mixed and placed

inside the pipe. End caps are attached and then the pipe is spun, first at

115 rpm to distribute the material and then at 1700 rpm for 2 minutes to

compact the PC (Figure 5). One end cap is removed and any excess slurry is

removed. A prelilminary cure at ambient temperature is accomplished with

he end caps in place. The pipe is then placed in a hot air circulating

oven (Figure 6) at 190°C and a final cure is attained in ~8 hours. The PC

liner can be seen in Figure 7

.

P40-6
D. Materials Properties Testing

Since the steel shell of the pipe can be expected to carry all or most

of the mechanical stresses, high physical strength in the PC are not of

primary concern. There are several properties that PC must have to provide

protection for the carbon steel shell. These include low permeability,

good bond strength to the steel inner surface, good corrosion and acid

resistance, and low erosion rates.

1. Permeability . Since the purpose of the liner is to provide a barrier between the geothermal fluid carried in the pipe and the mild steel

shell, the permeability coefficient of the lining material is important.

Studies were made on cylindrical samples 10 mm thick in the apparatus shown

in Figure 8. An average permeability coefficient was calculated, using

Darcy's Law, to be 22xl0“12 cm/sec. This is about 23 times less permeable than portand cement concrete.

Sections of 15.2-cm-i.d. pipe with a nominal 2.54-cm wall thickness

have been cast in PC with formulations previously mentioned. Specimens

have maintained fluid pressures as high as 2.75 MPa before bursting or showing any signs of leakage. & Attempts at measuring the permeability

coefficient of centrifugally cast PC pipes have been unsuccessful because

of leakage at the gaskets between the pipe ends and the fixture (Figure 9).

2. Bond Strengths . Bond strength between PC liner and the steel pipe

has also been tested. Bond strengths are measured on sections of PC-lined

steel pipe in a fixture as shown in Figure 10. Laboratory-cast sections

have shown bond strength to average ~5.5 MPa.

Sections of a PC-lined pipe that had been exposed to the 280,000-ppm

flowing brine at 235° to 250°C for 90 days at the Bureau of Mines Materials

Testing Facility at Niland, California, were tested at 25° , 100°, and

200°C. The bond strengths were 10.3, 10.4, and 8.9 MPa respectively. A

section of the pipe after the bond test is shown in Figure 11. The ability

of the PC liner to withstand the high force required to disbond it from the

P40-7
steel and the absence of salts or corrosion products on the steel are indications of the high strength and impermeability of the liner.

3. Corrosion and Acid Resistance . Testing of PC samples in the

laboratory and in the field in various geothermal environments has shown

the material able to withstand the most adverse conditions. Loss of

strength is generally noted for 30 to 60 days, then the materials stabilize

and exposures of up to 3 years in several geothermal environments indicate

good durability.

The feasibility of using PC materials in high temperature - low pH

environments has also been investigated. PC samples have been exposed to a pH 1 HC1 solution at 90°C for more than 440 days. No evidence of deterioration of the PC as determined by weight loss or volume change has been

detected. Additional tests have been conducted in an environment of pH 1 HC1 at 200°C. PC samples manufactured with several monomer systems with a silica sand - portland cement aggregate system have been exposed for 180

days with no apparent deterioration (Figure 12).

4. Erosion Resistance . No definitive studies have been conducted to

determine the erosion capabilities of PC. However, several PC-lined steel

pipes have been exposed to flowing brine for periods of up to three years

at the East Mesa geothermal site near El Centro, California. The brine flow

rate was 226-302 lit/rain at a temperature of 160°C. No deterioration,

erosion, or scale accumulation was detected upon completion of the test.

E. Economic Assessment

The Bechtel Corporation has performed conceptual designs for two 50-

MWe power plants operating on geothermal brines. ^ One plant utilized a moderate-temperature low-salinity brine as produced near Heber, California,

in the Imperial Valley, and the other utilized a high-temperature, highsalinity brine as produced in Niland, California, near the Salton Sea. In

both plants extensive use was made of carbon steel for vessels and

P40-8
piping, with Type 316 L stainless steel or titanium used only when carbon

steel was considered to be unsuitable.

Burns and Roe Industrial Services Corporation (BRISC) reviewed the two

designs with the view of substituting, where possible, PC materials for the

metals used by Bechtel.-*-® Each plant was considered separately and

estimates of cost savings which could be made in the capital cost of the

two plants were calculated. The capital cost estimates covered the battery

limits plants and the brine supply and reinjection costs as delineated in

the report for the Heber plant. No estimate of the piping was included for

the Niland plant because of the limited amovT't of detail provided in the

original report.

In the case of the Heber plant, it was shown that a savings of

~$880,000 can be achieved, based on the equipment and lines within the

battery limits, and ~$2,800,000 for brine supply and reinjection lines

outside the battery limits (Table 3) . In the Niland case a savings of

~$550,000 for equipment and lines within the battery limits was shown.

These savings in capital costs have a direct impact on the cost of

electric power and could result in a cost reduction of 2.72 mills per

kilowatt hour for the Heber plant. A similar savings is expected for the

Niland plant. It was also estimated that the use of PC-lined vessels and

piping would allow for an on-stream time of ~74% versus 70% using unlined

carbon steel components. This higher on-stream availability could result

in additional savings of ~3.46 mills per kilowatt hour. Thus a net savings

of ~6.2 mills per kilowatt hour could be realized by using PC materials

wherever possible.

CONCLUSIONS

1. The durability of PC in geothermal fluid environments has been

demonstrated by extensive long-term tests conducted at BNL and at six

geothermal sites.

P40-9
2. PC may be used in the design of geothermal or saline water facilities to counter corrosion at temperatures up to 250°C.

3. The cost of PC liners in most cases is less than other alternative

methods of providing corrosion protection.

4. PC, made with properly selected monomers and aggregates, can protect carbon steel surfaces at temperatures above the expected geothermal

operating range and preserve structural integrity without the use of corro- sion allowances.

5. The need for exotic materials such as titanium, molybdenum, or austenitic stainless steels (Type 316 L) is reduced or eliminated by using

PC-lined carbon steel.

6. Material costs for vessels and piping within the Heber plant

battery limits results in a 44% savings through the use of PC linings.

7. Brine supply and reinjection piping material costs can be reduced

by 40% using PC linings.

8. The use of PC materials at the Heber Plant can result in a reduction in the cost of electric power production by 2.72 mills per kilowatt

hour.

9. A 4% increase in on-stream availability for the Heber plant, due

to the use of PC liners, will result in a net cost reduction of ~3.46 mills

per kilowatt hour.

10. The reduction in capital cost and anticipated improvements in on- stream availability through the use of PC in the Heber plant will result in an overall saving of ~6.2 mills per kilowatt hout.

P40-10
References

1. Anderson, David N., and Lund, John W. , Editors, Direct Utilizat ion

of Geothermal Energy: A Technical Handbook , Geothermal Resources Council

Special Report No. 7, ISSN 0149-8991, ISBN 0-934412-07-03, 1979.

2. McCright , R.D. , Frey, W. F. , and Tardiff, G. E., Localized Corrosion of Steels in Geothermal Steam/ Brine Mixtures, Geothermal: Energy for

the Eighties, Transactions, Volume 4, Geothermal Resources Council, ISSN

10193-5933, ASBN 0-934412-54-5, Davis, California, 1990.

3. Schroeder, J.E., Design and Fabrication of Polymer Concrete Pipe

for Testing in Geothermal Energy Processes , Final Report, July 1981.

4. Kukacka, L. E., Fontana, J. , Auskern, A., Concrete Polymer

Materials for Geothermal Applilcations , Report No. 6, BNL 20571, July 1975.

5. Kukacka, L. E., et al., Alternate Materials of Construction for

Geothermal Applications, Progress Report No. 14, BNL 50751, September 1977.

6. Sugama, T. and Kukacka, L.E., Cement and Concrete Research, Vol.

9, pp. 69-76, 1979.

7. Kaeding, A., Design and Fabrication of Polymer Concrete-Lined Pipe

for Testing in Geothermal Energy Processes, Final Report, December 1981.

8. Kukacka, L.E., et al., Concrete-Polymer Materials for Geothermal

Applications, Progress Report No. 9, BNL 21665, June 1976.

9. Conceptual Design of Commercial 50 MWe (net) Geothermal Power

Plants at Heber and Niland, California. October 1976, Energy Research and

Development Administration, Division of Geothermal Energy, SAN-1124-1.

10. Economic Assessment of Polymer Concrete Usage in Geothermal Power

Plants, November 1977, Prepared by Burns and Roe Industrial Services

Corporation, BNL 50777.

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List of Figures

Figure 1 - PC samples tested at the Geysers, Calif, in flowing dry steam

at ~235°C.

Figure 2 - PC lined pipe installed at the Bureau of Mines Corrosion

Facility in Niland, Calif. Flowing geothermal brine

>280,000 ppm dissolved solids at ~240°C.

Figure 3 - PC lined pipe installed at East Mesa Geothermal Site near

El Centro, Calif, flowing brine ~160°C.

Figure 4 - Sandblasting inner surface of steel pipe prior to application

of PC liner.

Figure 5 - Centrifugal casting of PC liner in steel shell.

Figure 6 - Curing of PC liner in hot air circulating oven.

Figure 7 - Centrifugal cast PC liner.

Figure 8 - Permeability fixture for flat plates.

Figure 9 - Permeability fixture for PC pipe.

Figure 10 - Fixture used to measure bond strength of PC liner to steel

sheel

Figure 11 - PC liner removed from steel pipe after testing at Bureau of

Mines Corrosion Facility.

Figure 12 - PC samples exposed to pH 1 HC1 at 200°C.

P40-15
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12

ELECTROLESS NICKEL AS A SUBSTITUTE FOR

CHROMIUM PLATING IN INDUSTRIAL APPLICATIONS

Ronald N. Duncan

ELNIC, Inc.

Nashville, Tennessee

In the last decade American industry has become almost

totally dependent upon foreign supplies for strategic

metals and minerals. While presently we import only

4 2 percent of the petroleum we use, more than 80 percent

of the chromium, cobalt, manganese, niobium, tantalum

and platinum required for industry is supplied by foreign

producers.

Of these strategic metals, chromium is probably the most

critical. Chromium is indispensable for the manufacture

of stainless and alloy steels and is essential to the

metal finishing industry. The United States presently

consumes 25 percent of the world's production of chromium

and has no domestic reserves. Worse yet, 98 percent of

the world's chromium supply is located in southern Africa.

Because of this area's political instability, a stoppage

or reduction of its production is very possible (1 , 2)

In the National Materials Advisory Board's study on

"Contingency Plans for Chromium Utilization" (3) , electroless nickel was cited as the only true substitute for

chromium. All other suggestions were techniques to

merely reduce, not eliminate, chromium usage. Electroless nickel was also described in this report as a probable improvement over chromium in many applications.

This paper describes this coating and its properties

and discusses how it can be used as a substitute for

industrial chromium coatings and for other more expensive

or less readily available materials.

AN AMORPHOUS MATERIAL

Electroless nickel is an alloy of nickel and phosphorus.

Those coatings used for functional (rather than decorative)

applications typically contain 10 to 11 percent phosphorus

dissolved in nickel and less than 0.05 percent other

impurities. Unlike chromium or other electrolytic

coatings, electroless nickel is completely amorphous.

P41-1
It has no crystal structure and contains no segregation

or separate phases. Its lack of long range order has

been confirmed by electron diffraction studies at

magnifications of 150 , 000 ( 4 ). An example of this

coating is the 75 pm (3 mil) thick deposit shown in

Figure 1. It is the lack of structure of this alloy

which produces its unusual properties and makes it well

suited for protection against corrosion, erosion and

wear. Some of the properties of electroless nickel are

compared to those of commercial hard chromium coatings

in Table 1.

Because of their purity and homogeneity, the internal

stress of high phosphorus electroless nickel coatings

is very low on most substrates. As shown by Figure 2(5),

on steel the coating is typically compres sively stressed

at 4 MPa (0.5 ksi)

.

This helps to ensure the deposits

integrity and the coatings performance.

The internal stress of commercial hard chromium coatings

is always very highly tensile, and usually exceeds 200 MPa

(30 ksi). This, combined with chromium's brittleness,

causes the coatings to be cracked and often produces

fissures through its thickness to the substrate. An

example of the cracks in hard chromium deposits is shown

in Figure 3.

A similar problem occurs with electroless nickel coatings

containing less than 10 percent phosphorus. As illustrated

by Figure 2, the internal stress level of low phosphorus

deposits is also high. This also causes these deposits

to be cracked and porous. The discontinuities present

in both chromium and low phosphorus, electroless nickel

coatings reduce their strength, ductility, and wear and

corrosion resistance.

UNIFORMITY PROVIDES MANY BENEFITS

Unlike electroplated coatings, electroless nickel is

applied without an electric current. Instead, the

coating is produced by autocataly tic chemical reduction.

Electroless nickel is plated onto a substrate by reducing

nickel ions to metallic nickel with sodium hypophosphite

.

This chemical process avoids the non-uniformity associated with most other metallic coatings.

The thickness of electroplated coatings like chromium will

vary significantly depending upon the part's configuration

and proximity to the anodes. The coating tends to buildup

on corners, edges and the like, and to be reduced on

internal surfaces. Not only do these variations effect

P41-2
the ultimate performance of the coating, but they can also

cause additional finishing or machining to be needed after

plating. With electroless nickel, the coating thickness

is the same on any area of the part that is exposed to

fresh plating solution. Grooves, slots, blind holes, and

even the inside of tubing will have the same amount of

coating as the outside surface of the part.

The benefits of electroless nickel's uniformity are illustrated by its substitution for chromium on many of the

cylinders and rolls used in the printing and textile

industries. This change has reduced chromium usage, and

has also significantly reduced the finishing costs of these

components. Previously the cylinders had to be ground,

plated to a thickness of about 250 ym (12 mils) , and ground

a second time, before they could be balanced and installed.

With electroless nickel they are now ground only once,

balanced, plated to the desired diameter -- usually with

25 to 37 ym (1 to 1% mil) -- and installed. Not only has

this reduced the cost of plating by 40 percent, but 55

percent of the grinding time is also saved, freeing the

machines for the production of new parts and increasing

productivity (6)

The thickness of electroless nickel coatings can be

controlled to suit the application. Coatings as thin as

2% ym (0.1 mil) are commonly applied to electronic components, while those as thick as 75 ym (3 mils) are typical

for chemical or petroleum equipment. Coatings thicker

than 250 ym (10 mils) are also easily applied, but because

of cost are normally used only for salvage or repair of

worn or mis-machined parts.

GLASS-LIKE PROPERTIES

The mechanical and physical properties of electroless

nickel deposits resemble those of other glasses. They

have high strength, limited ductility and relatively

low conductivity. The ultimate tensile strength of high

phosphorus deposits exceeds 700 MPa (100 ksi)

This is

equal to many alloy steels and 5 to 10 times higher than

chromium. This allows the coating to withstand considerable abuse without damage. Accordingly, it is commonly

used to protect compressor blades, turbines, valves, pumps,

extruders, blowers, and the like, and has often replaced

stainless steel and exotic alloys.

The ductility of electroless nickel is only about 1 to 1%

percent. While this is less than that of most alloys, it

is adequate for most coating applications and much superior

to chromium. Thin films of the deposit can be bent

P41-3
completely upon themselves without fracture, and the coating

has been used successfully for springs and bellows.

Electroless nickel, however, should not be applied to

articles which subsequently will be bent or drawn. Severe

deformation will crack the deposit, reducing corrosion and

abrasion resistance.

The electrical and thermal conductivity of electroless

nickel is low. Its conductivity can be increased by heat

treatment, but is still much less than that of conventional

conductors like copper or silver. Because of the relatively

thin layers used, however, for most applications its resistance is not significant. Electroless nickel coatings are

being successfully used for such applications as exchanger

tubing and electrical switches and contacts.

Electroless nickel can be easily soldered, braised, and

bonded and is often used to ease joining of non-metals

and aluminum and stainless steel. This, combined with its

uniformity and corrosion resistance, have made electroless

nickel an ideal coating for electronic components. Accordingly, it is being increasingly used to reduce or eliminate

gold and precious metal requirements in the electronic

industry ( / )

.

PROVIDES EXCELLENT RESISTANCE TO WEAR

One of the most important properties of electroless nickel

for many industrial applications is its hardness and wear

resistance. As deposited, high phosphorus coatings have

a microhardness of 480 to 500 VHNl00» approximately equal

to 48 HRC. This is similar to many hardened steels.

Heat treatments, similar to age hardening procedures for

aluminum alloys, can produce significant increases in

coating hardness. As shown by Figure 4, hardness values as

high as 1100 VHN100 (approximately 69 HRC) can be produced.

This is equal to the hardness of commercial hard chromium

and comparable to that of some hard facing alloys and

ceramics

.

Hardening of electroless nickel is due primarily to the

formation of nickel phosphide particles within the alloy.

At temperatures above 260°C (500°F) coherent and then

distinct particles of Ni3P begin to form, and at temperatures above 320°C (600°F) the glass begins to crystallize.

This causes its hardness and wear resistance to increase

rapidly. Maximum hardening is obtained through treatments

at 400°C (750°F)

.

P41-4
Both heat treated and non-heat treated electroless nickel

coatings are commonly used to minimize the effects of

erosion, abrasion and wear. Laboratory tests have shown

fully hardened coatings to have wear resistance equal to

hard chromium. This is illustrated by Table 2, which

shows the results of Taber Abraser Wear tests of electroless nickel coatings, and compares them to electroplated

nickel and chromium(o). The excellent resistance of

electroless nickel often allows it to replace high alloy

materials, hard chromium, and even hard facings.

An example of the ability of electroless nickel to make

common materials behave like superior ones is the polyethylene pelletizer bowl shown in Figure 5. This bowl,

with a 125 pm (5 mil) thick high phosphorus coating,

was in service for two months with no measurable loss

or attack. Previous uncoated aluminum bowls failed in

less than three weeks after 40 percent of their weight

was eroded away. Other typical wear applications

include feed screws and extruders, computer drive

mechanisms, textile and fiber equipment, hydraulic

cylinders, molds and dies, and packaging equipment.

Another property related to wear is lubricity. Under

sliding or abrasive conditions, low friction surfaces

minimize heat buildup and experience less scoring and

galling than higher friction surfaces. Because of the

phosphorus they contain, electroless nickel deposits

have a low coefficient of friction, typically 0.13

(lubricated) to 0.4 (unlubricated). This is approximately 20 percent lower than chromium, one-half of that

of steel, and significantly better than aluminum or

stainless steel.

One disadvantage of electroless nickel is that, unlike

chromium, it is not hydrophilic. Electroless nickel

lacks the crack pattern present in chromium which can

hold an oil or water film. Thus for applications like

ink rollers, which depend upon surface wetting, electroless nickel may not always be suitable.

ADHESION STRENGTH IS OUTSTANDING

The performance of a coating is often dependent upon its

adhesion to its substrate. Without adequate adhesion

even the most resistant coating can become dislodged or

broken, allowing attack of the underlying metal. The

adhesion of electroless nickel coatings to steel, aluminum,

copper and their alloys normally exceeds the sheer strength

of the substrate.

P41-5
The high bond strength of these coatings is due to the

ability of the plating solution to completely remove

microscopic contaminants from the substrate prior to

the deposition of the first nickel-phosphorus layer.

This allows the coating to develop both mechanical and

metallic bonds with the substrate metal. The .adhesion

of electroless nickel coatings to steel and aluminum is

typically 300 to 400 MPa (40 to 60 ksi). This is

similar to and often superior to that of chromium coatings.

PROVIDES SUPERIOR CORROSION RESISTANCE

One of the most important differences between electroless

nickel and chromium coatings is their corrosion resistance. Both are barrier coatings. Both protect the

underlying metal by sealing it off from the environment

rather than by sacrificial action. Thus, to be completely

protective the coatings must be defect free.

Because of the cracks present through even thick deposits,

hard chromium coatings offer only limited protection

against corrosion. These cracks offer pathways through

the coating for a corrosive environment to reach and

attack the substrate. Accordingly, chromium coatings will

oftentimes develop a network of rust spots or rust lines

across its surface. In addition, unlike electroless

nickel, chromium coatings are subject to interface or

underdeposit attack. With chromium, corrosion not only

will extend into the substrate at the bottom of cracks

or pores, but will also travel out from these areas along

the coating- to-substrate interface, loosening and lifting

the coating(9).

Because of its homogeneity and freedom from defects,

electroless nickel coatings provide a true barrier to

corrosion. They do not offer any pathways to the substrate.

The metallic bonds these coatings form with their substrate

also prevent underdeposit attack. Even if a pore were to

be produced by improper processing or mechanical damage,

in most environments it rapidly fills with corrosion

product, stifling further attack. Corrosion will not

spread out from the defect and thus it is contained.

High phosphorus electroless nickel coatings are almost

totally resistant to alkalis, like caustics and potash;

to salt solutions and brines, like sea water or those

present in food or chemical environments; to acid gas

environments, like those in the petroleum industry; and

to all types of organic media and solvents. The coating

also has good resistance to ammonia solutions; to organic

P41-6
acids, like lactic or acetic; and to reducing acids,

like hydrochloric or sulfuric. It is only significantly attacked by str n * ‘\utions like

In most environments, the corrosion resistance of hard

chromium is much less than that of electroless nickel.

Chromium is rapidly attacked by reducing environments

and is subject to pitting and localized attack in

halogens, especially oxidizing halogens like ferric

chloride. This is illustrated by Table 3, which compares

the corrosion of electroplated and cast chromium to that

of electroless nickel in different environments ( 9 , 10 , 11 )

The petroleum and chemical process industries are the

largest users of electroless nickel for corrosion

protection. Because of its superior resistance to attack

in these environments, the coating ensures easy and

reliable operation and extends the equipment life.

Accordingly, it is often used in place of more critical

or expensive materials, especially stainless steel. For

instance, oil field valves coated with 75 ym (3 mils) of

electroless nickel cost approximately one-third that of

equivalent stainless steel valves and in most environments provide equal protection(12)

Many petroleum components now coated with electroless

nickel were originally plated with hard chromium. The

poor performance of chromium in oil field environments,

however, lead to a reduction in its use and to the

increased specification of electroless nickel(13).

Specific applications in these industries include ball,

gate, plug and check valves, blow out preventers,

chokes, heat exchange equipment, pumps, compressors,

tubing, vessels, packers, and all types of down hole

equipment ( 12)

Electroless nickel has many unique properties which make

it a superior engineering material. The coating offers

high strength, excellent abrasion and wear resistance,

lubricity, solderability , and superior corrosion resistance, together with ease of application and uniform

thickness. Accordingly, electroless nickel has proved

to be useful in improving reliability, in reducing cost,

and in avoiding the use of critical or strategic

materials

concentrated sulfuric

CONCLUSION

RND: ss

October 4, 1982

P41-7
REFERENCES

1. Gray, A.G., Metal Progress, Vol. 117. No. 3.

p. 33 (1980).

2. Swinburn, J., Materials Performance, Vol. 21,

No. 1, p. 54 (T9~8"2‘) .

3. National Materials Advisory Board, "Contingency

Plans for Chromium Utilization", National Academy

of Sciences, Washington, D.C., 1978.

4. Weil, R. , Stevens Institute of Technology, private

communication, January 24, 1980.

5. Parker, K. and Shah, H. , Plating, Vol. 58, No. 3,

p. 230 (1971).

6. Paliotta, J.V., "Functional and Economic Impact

of Electroless Nickel on the Printing Press

Industry", Electroless Nickel Conference II,

Cincinnati, March, 1981.

7. Bandrand, D.W.

,

Plating and Surface Finishing,

Vol. 68, No. 12, p. 57 (1981).

8. Industrial Nickel Plating and Coating, International

Nickel Company, New York, 1976.

9. Uhlig, H.H.

,

Editor, Corrosion Handbook , John Wiley

and Sons, New York, 1948, p"l 825-828

.

10. Duncan, R.N.

,

"Corrosion Control With Electroless

Nickel Coatings”, AES Electroless Plating Symposium,

American Electroplating Society, St. Louis, March,

1982.

11. LaQue , F.L., and Copson, H.R.

,

Corrosion Resistance

of Metals and Alloys , Reinhold Publishing, New

York , 1963 , p^ 448-449

.

12. Duncan, R.N., "Performance of Electroless Nickel

Coatings in Oil Field Environments", Corrosion/82

Conference, National Association of Corrosion

Engineers, Houston, March, 1982.

13. King, J.A.

,

and Badelek, P.S.C., Oil and Gas Journal,

Vol. 80, No. 28, p. 115 (1982).

P41-8
TABLE 1

COMPARISON OF ELECTROLESS NICKEL AND COMMERCIAL HARD CHROMIUM COATINGS

PROPERTY ELECTROLESS NICKEL COMMERCIAL HARD CHROME

MATERIAL Alloy of 10 to 11 percent

dissolved in nickel.

Chromium plus trace

amounts of oxides

and hydrogen.

STRUCTURE Amorphous; no phase structure,

lamination or segregation.

Crystalline; fine

grained with

numerous cracks.

INTERNAL STRESS

ON STEEL, MPa

<7 200-300

DENSITY, g/cm3 7.75 6.90-7.18

MELTING POINT,

°C

890 1610

ELECTRICAL

RESISTIVITY,

wfi-cm

90 14-66

THERMAL

CONDUCTIVITY

,

W/cm- °K

0.08 0.67

MAGNETIC COERCITY Non-magnetic Non-magnetic

TENSILE STRENGTH,

MPa

>700 <200

DUCTILITY, 1 ELONGATION

1 to 1% <<0.1

MODULAS OF

ELASTICITY, GPa

200 100-200

COEFFICIENT OF

THERMAL

EXPANSION,

ym/m/ °C

12 6

ADHESION STRENGTH,

MPa

300-400 Good

HARDNESS, VHN10 0 480 to 500, as deposited;

heat treatable to 1100

800 to 1000

COEFFICIENT OF

FRICTION VS

STEEL (LUBRICATED)

0.13 0.16

TABER WEAR

RESISTANCE, . mg/1000 cycles

15 to 20, as deposited;

2 to 9 after heat treatment

2 to 3

CORROSION

RESISTANCE

Excellent resistance to attack

by all but the most severely

oxidizing environments.

Poor due to cracks;

resists oxidizing

environments; attacked

by halogens and reducing

solutions.

RND: ss October 4, 1982

P41-9
TABLE 2

COMPARISON OF THE TABER ABRASER RESISTANCE

OF DIFFERENT ENGINEERING COATINGS .

COATING HEAT TREATMENT TWI , mg/ 1000 CYCLES (D

Watts Nickel None 25

Electroless

Nickel

None 17

Electroless

Nickel

300°C/1 hr 10

Electroless

Nickel

500°C/1 hr 6

Electroless

Nickel

650°C/1 hr 4

Hard Chromium None 2

(1) Taber Wear Index, CS-10 abraser wheels, 1000 gram

load, determined as average weight loss per 1000

cycles for total test of 6000 cycles.

RND: ss

October 4, 1982

P41-10
TABLE 5

COMPARISON OF THE CORROSION BEHAVIOR

OF CHROMIUM AND ELECTROLESS NICKEL

IN DIFFERENT ENVIRONMENTS

ENVIRONMENT CORROSION RATE, ym/y

CHROMIUM

ELECTROLESS

NICKEL

TEMPERATURE 12°cd) 16°C(2) 20 o C<3)

10% Acetic acid nil 660 25

10% Citric acid nil — 19

Cone. Hydrochloric acid -- 100,000 46

10% Hydrofluoric acid 25,000 -- 30

10% Lactic acid nil -- 19

10% Malic acid 51 -- 17

10% Nitric acid nil -- 44

Cone. Nitric acid nil -- >25,000

10% Phosphoric acid 25 5 16

10% Sulfuric acid 280 -- 12

Cone. Sulfuric acid 760 -- 25

10% Sodium hydroxide nil -- nil

10% Ammonium chloride nil -- nil

10% Cupric chloride 380 -- 25

10% Cupric nitrate 51 -- 12

10% Ferric chloride nil -- 780

10% Sodium chloride nil 0.5

(1) Commercial hard chromium deposit; corrosion rates less than

25 ym/y reported as nil.

(2) Cast chromium metal.

(3) Electroless nickel containing 10%% phosphorus and less than

0.05% other elements.

RND : s s October 4, 1982

P41-11
FIGURE 1

Typical high phosphorus

electroless nickel

deposit on a mild steel

substrate. The coating

thickness is 75 ym

(3 mils) . Lighter

layer at top is a

copper overplate applied

for edge resolution.

400X magnification.

Nital etchant.

FIGURE 2

Effect of phosphorus

content on the internal

stress of electroless

nickel deposits on

steel. (1 ksi = 6.89 MPa)

P41-12
FIGURE 5

Typical cracks in a

commercial hard

chromium deposit.

1000X magnification.

Not etched.

FIGURE 4

Effect of different

one hour heat

treatments on the

hardness of electroless nickel containing

10%% phosphorus.

Temperature Degree* F

P41-13
FIGURE 5

Plated aluminum pelletizer bowl, after two months of

exposure to high velocity, polyethylene pellets and

Mississippi River water. Bowl is coated with 125 ym

of high phosphorus electroless nickel. Previously,

uncoated bowls failed in 3 weeks with weight losses

of 40 percent.

P41-14
DEVELOPMENT OF DUCTILE POLYCRYSTALLINE Ni 3 Al FOR

HIGH-TEMPERATURE APPLICATIONS*

C. T. Liu and C. C. Koch

Metals and Ceramics Division

Oak Ridge National Laboratory

Oak Ridge, Tennessee 37830

ABSTRACT

The nickel aluminide, Ni 3 Al, is strong and stable at

elevated temperatures; however, low ductility and brittle

fracture restrict its use for structural applications.

Microalloying has been employed to overcome the embrittlement

associated with grain-boundary separation. Among the dopants

added, boron has been found to be most effective in improving

fabricability and ductility of polycrystalline Ni 3 Al. Tensile elongations of >50% have been achieved by control of

boron concentration and thermomechanical treatment. The

B-doped Ni 3 Al shows excellent strength and oxidation

resistance at high temperatures. Unlike conventional alloys,

the yield strength of the aluminide increases rather than

decreases with increasing test temperature. The potential for

developing the aluminide as a substitute for Cr-containing

heat-resistant alloys is assessed.

INTRODUCTION

Austenitic stainless steels and superalloys are the common heatresistant materials for structural use at elevated temperatures. The

alloys typically contain substantial quantities (15~30 wt %) of chromium

for oxidation and corrosion resistance. Because of the strategic nature

of the chromium supply, there has been an increasing interest in development of aluminides as a substitute for Cr containing alloys.

JL Research sponsored by the Exploratory Studies Program, Oak Ridge

National Laboratory and Division of Materials Sciences, U. S. Department

of Energy under contract W-7405-eng-26 with the Union Carbide

Corporation.

P42-1
2

Nickel and iron aluminides are in a class of materials usually

referred to as intermetallic compounds. The aluminides are typically

strong, stable, and resistant to oxidation and corrosion 1 at elevated

temperatures. Unlike conventional alloys, the static strength of many

aluminides shows an increase rather than a decrease with increasing test

temperature. * 0 A major problem with using aluminides as structural

materials is their reported brittle fracture and low ductility, particularly at lower temperatures . ^

>

4 Because of the poor fabricability

and low fracture toughness, the aluminides have not been considered as serious candidates for structural engineering applications.

The nickel aluminide based on Ni 3 Al (y" ) is an important

strengthening constituent of commercial nickel-base superalloys. It has

been known for years that single crystals of M3 AI are quite

ductile, but its polycrystalline forms are extremely brittle. » The

brittleness of such poly-crystals is associated with weak grain boundaries that cause brittle intergranular fracture without much plastic

deformation within the grains. During the past 20 years, efforts have

been spent on improving the ductility of aluminides , 8 * 10 but significant

progress has been achieved only recently through microalloying

processes . 11 Microalloying involves adding a small amount of dopants

(usually <1 %) for controlling chemistry and cohesion of the grain boundaries. In this study, a microalloying approach was employed to alleviate the grain boundary embrittlement problem of Ni 3 Al. Emphasis will

be placed on the effect of boron additions on ductilization of Ni 3 Al

alloys

.

EXPERIMENTAL PROCEDURES

The Ni 3 Al alloys containing 26 and 24 at. % A1 (13.9 and 12.7 wt %

Al) were doped with a small amount of Ce (0.1 wt %) , Mn (2.1 wt %), and

B (0 to 0.1 wt %) . The alloys were prepared by arc melting and drop

casting, using pure metal elements and Ni—4 wt % B and Ni—4 wt % Ce

master alloys. There was no significant change in weight during the

alloy preparation. The alloy ingots (2.5 x 5.3 x 0.6 cm) were first

P42-2
3

homogenized at 1000-1200°C, and then fabricated to sheets by either hot

rolling at 1200°C or cold rolling at room temperature with intermediate

annealing at 1000°C.

The microstructure of Ni

3 Al alloys annealed for various times at

1000°C was examined metallographically . The specimens polished by standard techniques were etched in a solution of 20 parts of H2O, 20 parts

of HN03

10 parts of HF, 20 parts of H3PO4, 10 parts of acetic acid, and

10 parts of HC1. The crystal structures of Ni 3 Al alloys were determined

by x-ray diffraction using Cu radiation.

Sheet specimens with a gage section of 12.7 x 3.2 x 0.7 mm were used for tensile and creep tests. Tensile tests were conducted on an Instron testing machine at a crosshead speed of 150 mm/s. To perform

the tensile tests at elevated temperature, a water-cooled quartz-tube

vacuum system was attached to the Instron machine, and specimens were heated inductively inside a tantalum susceptor. Creep tests were done

in vacuum under dead-load arrangement. During the tests, the temperature was monitored by a Pt/Pt—10% Rh thermocouple centrally located

on the specimen. Fracture surfaces of selected tensile specimens were examined by a JSM-U3 scanning electron microscope (SEM) operated at

2 5 kV.

RESULTS

Undoped Ni3Al alloys containing 26 and 24 at. % A1 cracked badly

during hot and cold fabrication. Figure 1(a) shows part of the ingot

hot rolled at 1200°C. The ingot was almost pulverized due to extensive

cracking along the grain boundaries, indicating the grain-boundary

brittleness in Ni

3 Al. The alloys doped separately with 0.1 wt % Ce,

0.1% B, and 2.1% Mn all cracked during fabrication except the B-doped

Ni—24 at. % A1 alloy (12.7 wt. % Al), designated as IC-6. It was possible to fabricate the IC-6 ingot into sheet by repeated cold rolling

at room temperature and heat treatment at 1000°C. The amount of cold

work was initially -15% reduction in thickness, and was gradually

increased to 50% between each intermediate annealing. Figure 1(b) shows

the 0. 76-mm-thick sheet of IC-6 fabricated by cold rolling.

P42-3
4

(b)

Fig. 1. Comparison of fabricability of Ni

3 Al alloys

containing 24 at. % Al; (a) A part of Ni

3 Al ingot hot

rolled at 1200°C, showing extensive grain-boundary cracking; (b) A sheet of Ni

3 Al doped with 1000 ppm B, fabricated by cold rolling at room temperature.

P42-4
5

To study the effect of B additions on fabricability , a series of

alloys was prepared based on Ni-24 at. % Al, in which various levels of

B were added. Table 1 summarizes the results. The alloys containing

100 ppm B or less cannot be fabricated by cold rolling. IC-18 con- taining 250 ppm B was fabricated into sheets, but it exhibited numerous

shallow surface cracks. The alloys containing 400 ppm and above were fabricated into good quality sheets.

Table 1. Effect of boron additions on fabricabilitya of

Ni-24 at. % (Ni-12.7 wt %) Al

Alloy

B Addition

(ppm) Fabricability

IC-2 0 Cracked badly

IC-20 50 Cracked badly

IC-19 100 Cracked

IC-18 250 Sheet fabricated but exhibited

shallow surface cracks

IC-21 400 Sheet fabricated

IC-15 500 Sheet fabricated

IC-28 700 Sheet fabricated

IC-6 1000 Sheet fabricated

a Ingots were fabricated by repeated cold rolling

(15 ~ 50% reduction in thickness) and softening heat treatment at 1000°C.

The crystal structure in undoped and B-doped Ni 3 Al alloys was determined by x-ray diffraction. The superlattice lines, which were clearly visible, characterize the Ll 2 _type ordered crystal structure.

Thus, the small amount of B does not affect the long-range ordered

crystal structure in Ni 3 Al. The microstructure of B-doped Ni 3 Al alloys

(24 at. % Al) was examined as a function of aging time at 1000°C. The

alloys show a "regular" equiaxed grain structure with few annealing

twins and second-phase particles (Fig. 2). Grain growth occurs at

P42-5
6

Fig. 2. Microstructures of Ni 3 Al doped with 1000 ppm B

(a) 10 min anneal at 1000°C; (b) 16 d anneal at 1000 C.

P42-6
7

1000°C; however, there is no apparent difference in grain size for

specimens containing 500 and 1000 ppm B after the same annealing

treatment

Tensile properties of B-doped Ni 3 Al alloys were characterized as functions of B concentration, heat treatment, and test temperatures.

Table 2 shows the room-temperature properties of the alloys doped with

different levels of B. For the specimens annealed 30 min at 1000°C, the

yield strength showed a linear increase with the B concentration

[Fig. 3(a)] but the tensile strength was essentially insensitive to the

B concentration. All the B-doped alloys exhibited transgranular ductile

fracture [like Fig. 4(a)] with the ductility exceeding 40%. The tensile

elongation increased with the B content [Fig. 3(b)] and reached 53.8%

for the IC-6 specimen containing 1000 ppm B. After being annealed for

16 d at 1000°C, the IC-6 specimen showed a higher yield strength but

distinctly lower ductility. In accompanying the decrease in ductility,

the fracture mode changed from the transgranular to a mixed one [Fig.

4(b)]. On the other hand, the IC-15 specimen containing 500 ppm B

displayed a lower yield strength with ductility remaining unchanged

after the long term anneal.

Figure 5 is a plot of yield strength as a function of test temperature for the B-doped Ni 3 Al (IC-15) and commercial fabricable alloys

such as Hastelloy-X and type 316 stainless steel. Unlike the conventional alloys, the strength of IC-15 increases with increasing temperature and reaches a maximum around 600°C. As a result of this

increase, the aluminide is much stronger than the commercial alloys at

elevated temperatures. For example, the aluminide displays a yield

strength of 550 MPa (80,000 psi) and a ductility of 45% at 600°C. In

comparison, the yield strengths of Hastelloy-X and type 316 stainless

steel are 210 and 120 MPa (31,000 and 17,000 psi), respectively. Thus,

the aluminide has the superior high-temperature strength.

Table 3 shows the results of preliminary measurements of creep properties of the B-doped alloys tested at 760°C (1400°F) and 138 MPa

(20,000 psi). The minimum creep rate of the aluminides is insensitive

to the B concentration between 250 and 1000 ppm. IC-6 specimens exhibited a sharp drop in creep rate with an increase in grain size from

P42-7
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P42-8
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content

P42-9
10

Fig. 4. SEM fractographs of B-doped Ni 3 Al (1000 ppm B) tested in

tension at room temperature. (a) Specimen annealed for 3 d and 15 h at

1000°C. (b) Specimen annealed for 16 d at 1000°C.

P42-10
11

YIELD

STRENGTH

(MPa)

100

90

80

70

60 YIELD

STRENGTH

50 (ksi)

40

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20

10

0

Fig. 5. Yield strength as a function of test temperature for

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alloys Hastelloy X and type 316 stainless steel.

P42-11
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P42-12
13

25 to 120 um produced by annealing at 1000°C. The creep rate of B-doped

N1 3 A1 is substantially lower than that of type 316 stainless steel 12 and

Hastelloy X, 13 as indicated in Table 3.

An alloy coupon of IC-6 was exposed to air at 900° C for a study of

air oxidation. The sample was removed from the furnace after each

100-150 h exposure. The exposed sample displayed no apparent spalling

with a weight gain of 0.61 mg/cra2 after a total of 524-h exposure.

Metallographic examination indicates the formation of a thin oxide scale

[Fig. 6(a)] (— 2.6 ym thick) which is enriched with aluminum as revealed

by electron microprobe analysis [Fig. 6(b)]. Line scanning revealed a depletion of A1 in the base metal beneath the Al-rich oxide scale. A

strip cut from the oxidized coupon was bent to ~360 degrees without

failure (Fig. 7), indicating the B-doped aluminide remained ductile

after the exposure. All the results demonstrate the good oxidation

resistance of B-doped Ni^Al containing no Cr, the critical strategic

element

DISCUSSION

Since Ni 3 Al classically exhibits brittle intergranular fracture,

microalloying has been used to alleviate the grain-boundary embrittlement problem of the aluminide. In this study, two types of dopants were

added to Ni

3 Al. Type I dopants are Ce and Mn that, based on thermodynamic considerations, 14 have a strong affinity for impurities such as sulfur. Sulfur has been known to have a strong tendency to segregate to

grain boundaries in Ni-base alloys 15 as well as in Ni 3 Al 15 and

Ni 3 (A1 ,Ti)

16 The purpose of adding the dopants was to scavenge sulfur

and other harmful impurities from grain boundaries through a precipitation reaction. A Type II dopant is B, which tends to segregate at grain

boundaries and possibly increase the cohesive strength of the boundaries. Among the dopants we added, boron is found to be the most

effective in improving fabricabi lity and ductility of the nickel aluminide. This finding is in agreement with the results of Aoki and

Izumi 11 who first discovered the beneficial effect of B in Ni

3 Al. They

reported a tensile ductility of ~35% elongation at room temperature. By

P42-13
14

(a)

/3 M»)

Fig. 6. Formation of Al-rich oxide scale on the surface of the

B-doped Ni 3 Al (1000 ppm B) specimen exposed for 52A h in air at 900°C.

(a) oxide scale with an average thickness of 2.6 pm. (b) A1 mapping

obtained by electron microprobe analysis.

P42-14
15

Fig. 7. A bent strip of B-doped Ni 3Al (1000 ppm B) oxidized for

500 h at 900° C in air.

control of B concentration and thermomechanical treatment, we have

observed a tensile ductility of 53.8% elongation at room temperature —

the highest tensile ductility ever achieved by polycrystalline

aluminides. Using Auger electron spectroscopy. White 17 has detected a strong segregation of B to grain boundaries in B-doped Ni 3 Al. We

believe the presence of B at grain boundaries improves the cohesion of

the boundary and thus reduces the tendency toward the integranular fracture in Ni

3 Al. It should be pointed out that B additions only effectively ductilize the Ni 3 Al alloys containing 24 at. % Al, but not the

alloy containing 26 at. % Al. The reason for this behavior is not well

understood, but it is possibly related to the alloy stoichiometry, which

is known to affect the chemistry of grain boundaries and consequently

their segregation behavior. 18 * 19

The ductility and fabrication behavior of B-doped Ni 3 Al

(24 at. % Al) are dependent on the B concentration. The Ni 3 Al alloys

doped with <250 ppm B cannot be satisfactorily fabricated into sheets

P42-15
16

because of grain-boundary cracking. On the other hand, sheet fabrication becomes somewhat difficult for the alurainide doped with 1000 ppm.

Also, Ni 3 Al containing 1000 ppm B loses most of its ductility after

long-term annealing at 1000°C. Thus, the optiraiura level of B in Ni 3 Al

should be around 500 ± 200 ppm, based on the consideration of ductility

and fabricability

.

The yield strength of the B-doped Ni

3 Al (IC-15), just like undoped

Ni

3 Al, 2(1 shows an increase with temperature, and reaches a maximum

around 600°C. The positive temperature dependence of the yield

strength, which has been observed in other LI 2 ordered alloys, is

generally believed to be associated with a thermally activated hardening

process21 in connection with the {111} slip system. The strength starts

to decrease with temperature above 600°C, corresponding to a gradual

change in slip systems from {111} to {100} planes . (22, 23 ) The g-doped

Ni 3 Al alloys are much stronger than the existing commercial fabricable

heat-resistant alloys as shown in Fig. 5. Currently we are working to

increase the strength of Ni

3 Al further by both solid solution hardening

and particle strengthening.

The B-doped Ni

3 Al alloys are quite resistant to creep deformation

and air oxidation. The minimum creep rate of the fine-grained Ni

3 Al

specimens is significantly lower than that of type 316 stainless

steels 12 and Hastelloy X. 13 Furthermore, the creep rate of Ni

3 Al specimens can be substantially reduced (by a factor of 20) by an increase in

grain size from 25 to 120 pm. The higher creep rate in the fine-grained

specimens may be related to grain-boundary sliding which can be minimized by coarsening the grain structure. Metallographic examination

reveals numerous cracks formed along grain boundaries in these creepruptured specimens. The oxidation rate of B-doped Ni

3 Al (IC-6) is com- parable or lower than that of current heat resistant alloys 24 (depending on the Cr content in these alloys). The good oxidation resistance

results from the formation of aluminum oxide scales which protect the

alloy from excessive oxidation. Recent work2

5

in coating materials has

demonstrated that alumina-forming materials provide an excellent

resistance to hot corrosion in coal energy conversion systems. This

indicates the potential use of the aluminides as structural materials in

sulfiding environments.

P42-16
17

In summary, the aluminide N13A1 can be made ductile by

microalloying with a few hundred ppm boron. The B-doped Ni 3 Al alloys,

possessing a unique combination of high-temperature strength, oxidation

resistance, and adequate ductility and fabricability , have the potential

to be developed as new structural materials containing no critical strategic elements such as Cr and Co.

ACKNOWLEDGMENTS

The authors wish to express their sincere thanks to J. 0. Stiegler

and C. L. White for valuable technical discussions. Thanks are also due

to R. L. Heestand and L. Queener for alloy preparation and fabrication,

R. E. Clausing for electron microprobe analysis, 0. B. Cavin for x-ray

diffraction, W. H. Farmer for metallography, and E. L. Lee and

J. F. Newsome for technical assistance. The authors are also grateful

to Connie Dowker for preparation of the manuscript.

REFERENCES

1. E. A. Aitken, Intermetallio CompoundsJ. H. Westbrook, ed.,

pp. 491—515, John Wiley & Sons, Inc., New York, 1967.

2 N. S. Stoloff and R. G. Davis, Prog . Mater . Sci. 13(1), 1—34

(1966).

3. B. H. Kear, C. T. Sims, N. S. Stoloff, and J. H. Westbrook, eds.

"Ordered Alloys - Structural Applications and Physical Metallurgy,

Proa. 3rd Bolton Landing Conf. t September 1969, Baton Rouge,

Claitor’s Publishing Division, 1970.

4. H. A. Lipsitt, D. Shechtman, and R. E. Schafrik, Met . Trans. A

11A, 1369 (1980).

5. S. M. Copley and B. H. Kear, Trans . AIME 239, 977 (1967).

6. R. Moskovic, J. Mater. Sci. 13, 1901 (1978).

P42-17
18

7. K. Aoki and 0. Izumi , Trans. Jpn. Inst. Met. 19, 203 (1978).-

8. E. M. Grala, Mechanical Properties of Intermetallic Compounds,

J. H. Westbook, ed . , p. 358, John Wiley & Sons, Inc., New York,

1960.

9. A. V. Seybolt and J. H. Westbrook, Acta Metall. 12, 449 (1964).

10. M. J. Blackburn, et al., AFML-TR-78~18 and 78, Air Force Materials

Lab., Wright-Patterson Air Force Base, Ohio, 1977—78.

11. K. Aoki and 0. Izumi, Nippon Kinzoku Gakkaishi , 43, 1190 (1979).

12. V. K. Sikka, et al., ORNL-5285, Oak Ridge National Laboratory,

October 1977.

13. C. R. Brinkman, et al., 0RNL/TM-5405, Oak Ridge National Laboratory,

October 1976.

14. 0. Kubaschewski and C. B. Alcock, Metallurgical Thermochemistry 5th

ed. , Pergamon, New York, 1979.

15. W. C. Johnson, J. E. Doherty, B. H. Rear, and A. F. Giamei, Scr.

Metall. 8, 971 (1974).

16. C. L. White and D. F. Stein, Metall. Trans. A. 9A, 13 (1978).

17. C. L. White, unpublished results, Oak Ridge National Laboratory,

1982.

18. A. V. Seybolt and J. H. Westbrook, Acta Metall, 12, 449 (1964).

19. J. H. Westbrook and D. L. Wood, J. Inst. Met. 91, 174 (1963).

20. 0. Noguthi, Y, Oya, and T. Suzuki, Metall. Trans. A 12A, 1647

(1981).

21. S. Takenchi and E. Kuramoto, Acta Metall. 21, 415 (1973).

22. R. G. Davies and N. S. Stoloff, Trans. Metall. Soc. AIME 233, 714

(1965).

P42-18
19

23. C. Lall, S. Chin, and D. P. Pope, Metall. Trans . A 10A, 1323 (1979).

24. J. Stringer and I. G. Wright, Oxid. Met . 5, 59 (1972).

25. P. R. Clark, C. M. Packer and R. A. Perkins, FE-2592-29, Lockheed

Palo Alto Research Laboratory, March 15, 1981.

P42-19
t
INTRODUCTION

During the past 12 years, the interest in ceramics as hightemperature structural material has grown in the major industrial countries.

Through various governmental agencies, each country has channelled substantial

funds for producing and evaluating critical ceramic components for such

applications. Around 1970, the Advanced Research Projects Agency (ARPA) of

the U.S. Department of Defense, initiated a program to demonstrate successful

use of ceramics in demanding high-temperature structural applications.

Components for the hot flow path of a vehicular gas turbine engine

and some diesel engine parts are being extensively investigated. Of the

ceramic materials, dense silicon nitride and dense silicon carbide are considered the favorite candidates. Lately, toughened zirconia has been con- sidered favorably for some specific applications.

Shapes for most of these component parts are extremely complex.

Furthermore, the components must meet strict dimension tolerances and maintain

very high reliability in use. Even so, it is a prerequisite to produce these

components in large numbers at a competitive cost. Several fabrication

techniques can be used to produce such complex components. Many of these

techniques are being explored. The complexity of the components and the need

for large quantities dictated the consideration of injection molding as a technologically and economically feasible fabrication approach.

As a result of worldwide interest in this technology, Battelle

offered an industrial group research program to determine the important

factors in the ceramic injection molding process and in molding advanced

ceramics. Due to proprietary nature of our study, we have kept this paper

fairly broad and used information available in the open literature. Still,

the paper contains sufficient detail to underscore the various technical

aspects of ceramic injection molding.

COMPONENTS OF CURRENT INTEREST

Current efforts are focused on fabricating the following four

components:

• Turbine Rotors: To operate at an equivalent turbine inlet temperature of 1200 to 1400 C over the complete speed range (Figure

la).

P43-1
2

• Turbine Stators: To operate at 1400 C turbine inlet temperature

and withstand prevailing thermal and mechanical stresses during

startup, steady-state, and shutdown conditions (Figure lb).

• Supercharger Rotors: To replace expensive nickel alloys by lower

cost, nonstrategic materials. To investigate the possibility of

developing a 3-shaft regenerative turbine operating at 1300 C. It is believed that a 3-shaft turbine may have a better partial

load economy, two stages of compression would allow intercooling

between compressor stages and improve the specific fuel consumption at high load factors and also the specific power output;

also, with two stages of compression, relatively small turbines

driving the compressors can be designed to operate at modest

rotational speed and with reduced centrifugal stresses

(Figure lc).

• Combustion Prechamber Inserts: To replace high-nickel alloy in

diesel engine combustor components. Also to reduce ignition

delay on cold start up and to reduce idling/light-load noise

(Figure Id).

INJECTION MOLDING PROCESS

For some years, ceramic manufacturers have been aware of the

potential for injection molding; nevertheless, acceptance of this process as a routine production approach has been extremely slow. This skepticism can be

better understood if one presents the process technology in some perspective.

Ceramic injection molding consists, essentially, of the following

steps:

• Tailoring the ceramic powder

• Developing some organic binder formulations

• Producing a homogeneous ceramic and organic binder mix

• Forming parts by injection molding

• Processing the parts to remove the organic binder and densifying

the parts.

P43-2
3

Injection molding is obviously not a direct analog of the plastic molding

process. Successful ceramic injection molding depends on several process

steps. In this paper, we briefly discuss all these process steps and include

some experimental data. A silicon nitride powder has been chosen as an example.

Tailoring the Powder

Trend in ceramic industries is to use powders with high surface

reactivity to reduce the sintering temperature and time, and to achieve

improved end properties. GTE Sylvania's SN502 (characterized in Table 1) is one of the few appropriate commercially available silicon nitride powders. It

is typically composed of a mixture of amorphous, alpha and beta phases with

the alpha phase predominant. Its major contaminants are oxygen and free

silicon. The SEM micrograph in Figure 2 indicates that this powder is com- posed of a variety of shapes and sizes with a large fraction of submicron-size

particles and some agglomerates. Because the powder is a mixture of both

needle shaped (a-Si' 3N4 ) and equiaxed particles, a reliable size distribution can not be achieved. The needles vary from 0.1 to 7 or 8 microns in

diameter and from submicron to 20 to 30 microns in length, although the

majority are much shorter. The surface area indicates a large fraction of

particles less than 1 micron. The mixture of shapes in this powder drastically affects its packing and flow characteristics. One can easily understand

this in considering how a pile of needles or pins pack. Needle shapes persist

throughout the compaction process and lead to a relatively low compact

density. Low compact density causes excessive shrinkage during sintering.

High and nonuniform shrinkage becomes very critical for complex parts with

tight tolerances. Furthermore, such powders create problems when mixed with

relatively small amounts of organic binder, thus affecting the rheology of the

molding composition.

Moreover, covalent silicon nitride is not sinterable without an added sintering aid. In recent years, extensive research has been undertaken

to find chemical additives that would help the densif ication process.

Currently, Y 2 O3 or a mixture of Y 2 O3 and A1 2O 3 are widely used as a suitable

P43-3
4

aid for pressureless sintering of silicon nitride. Homogeneous mixing of the

additive(s) with silicon nitride powder is extremely critical. Thus, a powder

processing step is essential with SN502 silicon nitride to improve its flow

and packing and to render homogeneous dispersion of the additive(s). The

powder processing in this context refers to one of several methods of milling,

such as ball milling, fluid energy milling, or attrition milling. For silicon

nitride, dry ball milling has received the most attention and has been used as

a standard practice. The milling conditions and results are given in Table 2.

Irrespective of the type of milling aid, compaction density increased significantly after 16 hours. However, the breakdown of agglomerates and needles

became significant only after 24 hours of milling, as shown in the micrograph

in Figure 3 for the powder milled with stearic acid. This improvement represents deagglomeration and reduction in needle size. It is not necessarily a result of real particle size reduction during milling process. This powder

was used in a subsequent injection molding process.

Developing a Binder Formulation

Of primary importance is the selection of a binder formulation that

will give correct flow properties to a ceramic and binder combination such

that a void-free, stress-free, and uniformly dense molded shape can be

achieved. The flow behavior of such materials (shown in Figure 4 from

Metzner, 1959) can be described as:

• St. Venant where the viscosity is inversely proportional to shear

rate

• Bingham plastic which does not show any flow unless the shear

stress reaches a critical unit. The nonlinearity has been

attributed to changes in the network structure and alignment

within the network structure of the flowing material as the

stresses change. The network structure can support a stress

before any material flow can be initiated. This finite stress is

known as the yield stress.

P43-4
5

• Pseudoplastic where the viscosity decreases continually with

shear rate.

• Dilatant where the viscosity increases with shear rate.

Ideally, in injection molding, the mix leaving the gate of the mold cavity

should spread like a liquid sweeping the air before it, the air escaping

through mold clearances or vents. However, the absence of a yield point would

mean that the molded part may deform at the dewaxing stage. On the other

hand, if the mix has a high yield point, it will extrude into the mold cavity

as a thread or ribbon which coils upon itself, tending to trap air and induce

mechanical stresses as shown in Figure 5. To obtain best results, it is

necessary to select a binder formulation that will have a yield point con- sistent with the shape, size, and dimensional stability of the molded part.

In other words, the ceramic/polymer or wax formulations should have either

Bingham or pseudoplastic flow characteristic. Furthermore, the binder should

be easily degradable or extractable in a subsequent dewaxing process.

A binder formulation that will meet the criteria mentioned above has

been sought since 1932. An impressive list of patents that can be produced

covers almost all the natural and synthetic organics available in commercial

quantities. In general, low-molecular-weight polyethylene, polystyrene, polypropylene, and, more recently, paraffin and microcrystalline waxes with small

additions of other organic modifiers have been used successfully as binders

for ceramic injection molding. In the light of flow behavior and safe

dewaxing, we have chosen a molecular weight (^5000) polyethylene with a melt

index of 250.

Rheological Behavior . Experience has shown that good flow during

molding requires a viscosity of less than 10^ poise. As the shear rate often

ranges between 100 sec~l and 1000 sec

- *, it is assumed that any binder

formulation with a viscosity less than 10^ poise within the shear rate range

indicated will be suitable for molding. An Instron Capillary Rheometer which

has a shear rate range from 1 to 10^ sec

-

* has been used to describe the flow

behavior. Figure 6a shows the viscosity shear rate curves for unfilled polyethylene and polyethylene containing silicon nitride at different levels of

loading. Obviously, the addition of silicon nitride has significantly

P43-5
6

increased the viscosity at each shear rate but not changed the general shape

of the curve. At low shear rates, the viscosity of this formulation rises

sharply indicating a yield stress below which the material does not flow. It

also shows that the viscosity is outside the molding range at silicon nitride

loading of only 48 volume percent. Therefore, the viscosity of this formulation should be reduced so as to increase the silicon nitride loading, thereby,

increasing the density of the part and reducing the shrinkage.

The viscosity of such a formulation can be reduced by introducing

some organic additives in small concentrations. These are broadly known as modifying agents. Modifying agents should be compatible with the polyethylene

and should easily wet the silicon nitride powder surface. Paraffin, microcrystalline waxes, stearates, oleates, and many saturated and unsaturated

fatty acids can be used as modifying agents. Figure 6b illustrates the effect

of modifying agents on the silicon nitride loading and viscosity. By adding

18 volume percent of modifying agent the viscosity of the binder formulation

has been significantly reduced, even with an increased silicon nitride loading

of 60 volume percent. Obviously, the role of these modifying agents is rather

complex and some studies are in progress to develop some basic understanding.

Producing a Homogeneous Mix

A workable binder formulation thus developed can be used to produce

a sizable quantity of materials for molding. Production of such mix requires

an effective mixing technique. The following factors should be taken into

account when selecting a mixing technique:

• Knowing and defining where a homogeneous mix is attained

• Selecting a mixing system that moves the ingredients to be mixed

into a pattern that fulfills the requirements of the desired mix.

In mixing viscous materials, the major mechanism is shear. Shear

performs mixing by drawing out the components over thinner layers, which

reduces the size of the region occupied exclusively by one component. Various

types of commercial shear mixers are available. The Brabender Plastograph

with a sigma mixing head is an efficient mixing system.

P43-6
7

Forming the Parts

Basically, injection molding is a simple cyclic process in which the

granular mix is heated until softened, then forced into a mold where it cools

and resolidifies to produce a part in the desired shape.

Molding Machines . Many types and sizes of machines are available,

however, schematics of two basic types are shown in Figure 7. The screw-type

injection molding machine differs from a plunger type in material conveying

and plasticizing. A plunger machine heats the material by conduction and

convection only. In a screw machine, a major portion of heat comes from the

frictional forces between screw, material, and cylinder. In a plunger

machine, the pressure loss in the cylinder could be significant, and the

pressure at the end of the plunger is much higher than the pressure at the

nozzle. In a screw machine, the ram pressure and nozzle pressure are almost

the same. Selection of a specific machine type is influenced by the requirements of the molded product and by the degree of versatility desired. Regardless of the type, the following features in an injection molding machine are

of absolute necessity.

• Injection cylinder should have three individually controlled

heater zones. (Nozzle temperature should be separately

controlled.)

• Injection and follow-up pressures should be variable up to at

least 14 MPa, and be controlled by individual timers.

• Variable speed ram should be controlled either by a flow control

valve or by a timed hydraulic pump.

• Feed adjustment should be precise.

• Cycle timers should control plunger-forward time, mold-close

time, etc.

• Preposition and multiple prepack controls should be on the

plunger-type machine.

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8

Molding Cycle . A molding cycle is based on time, temperature, and

pressure. Being a dynamic process, its steps usually overlap; however, they

need to be characterized as identifable increments if analysis and control are

to be attempted. The cycle time is best represented by the diagram shown in

Figure 8 (from Gilmore and Spencer, 1950). The sequence is as follows:

t Period A (dead period) is the time (after the mold closes and the

screw/plunger starts forward) before material starts to flow into

the sprue bushing.

§ Period B (mold filling period) is the time during which material

flows into the mold and fills the cavity.

t Period C (pressure build-up period) is the time (after the cavity

is filled and the material begins to cool) during which pressure

in the cavity builds up to injection pressure.

• Period D (packing period) is the time (including Period C) during

which material flows into the cavity at a relatively low rate.

The injection cycle ends with this period and the screw/plunger

retracts.

t Period E (discharge period) is the time during which material in

the gate is still fluid and, with zero pressure on the runner

side and slightly less than hold pressure in the cavity, the

material flow reverses. (Flow rate is determined by the seal or follow-up pressure).

• Period F (cooling period) is the time during which material in

the cavity cools with no flow through the gate in either

direction.

• At the end of the mold closed time, the mold opens with some pressure remaining in the cavity. This pressure has been termed

the residual mold pressure.

Along with time sequence, it is essential to have the correct

temperature and pressure sequence. The pressure-temperature functions in

molding have been presented as "ideal gas" equation

(P + it) (V - W)=RT,

where P is the hydrostatic pressure, V is the specific volume, T is the

material temperature, and tt, W, and R are constants for specific molding

P43-8
9

material. It can be argued, if the melt temperature is constant in the cavity

immediately upon injection, the effective specific volume (V-W) is inversely

proportional to the effective applied pressure ( P+tt) . Once the gate is

frozen, the effective specific volume (V-W) is constant, and the effective

applied pressure (P+tt) drops in proportion to the temperature. It is possible

to construct a model from this equation and obtain some temperatures and

pressures for a specific material being molded into a fairly simple shape.

In the final analysis, molding success depends on the balance

between these parameters, the mold design, and the rheology of the molding

compound.

Mold Design . The success of injection molding may be measured by

the efficiency of the mold design and the quality of the mold construction.

Initial consideration should be given to parts design before designing a mold.

Feasibility of molding a part by injection molding should be based on such

facts as size and weight of the part, section thickness, shrinkage,

tolerances, draft, threads, radii, and holes. After careful review of all

information on a part, the actual mold design should be considered in the

light of the following factors:

• Runner length should be short; but diameter can be varied.

t Preferred runner shape is circular.

• Runner should be placed in the ejector half.

• Preferred gate shape is either round or rectangular.

• Gate size rather than shape appears to have more influence on the

manner in which the cavity is filled.

• Larger gate sizes appear beneficial because: — As the gate size is increased, the probability of jetting is

reduced by the lower velocity of the material entering the

mold — The evidence of knit lines is reduced — Cavity filling is more rapid and there is less opportunity for

the material to freeze prior to complete filling of cavity.

P43-9
10

• Changes in gate location can reduce or eliminate jetting of the

material

.

• Vent location in mold is most effective when jetting is reduced

and the cavity is filled from the gate area through the remainder

of the cavity. (Jetting may cause covering of the vent during the

initial stages of cavity filling.)

• Filling is more controlled when the material comes into contact

with a surface of the cavity immediately upon entry.

• Mix being injected into the cavity takes the path of least re- sistance; thus, using gates of different sizes simultaneously may

not be advisable.

• Depending upon the freezing rate of the material, it may not be

undesirable to use cavities of different sizes simultaneously.

• Rheological characteristics of the material influence the

retention of weld marks and knit lines in the molded parts.

• Wear on the machine parts and mold is still an unknown factor.

Removing the Binder

The binder must be removed (dewaxed) from the parts prior to

sintering. Binder removal is a key operation in the success of injection

molding. In the experimental design, the following factors should be taken

into account:

• The binder systems may be composed of more than one organic

ingredient with markedly different melt viscosity and

decomposition ranges. Hence, in such systems, attempts should be

made to remove the ingredients slowly and gradually to create

passages within the part with increasing temperature and time,

and to allow the major binder ingredient to escape at high

temperatures without causing any failure.

• Thermal degradation of nonoxide injection molded ceramics in air

could cause a serious oxidation problem. Under these conditions,

silicon carbide is more likely to suffer from oxidation than

silicon nitride. With silicon nitride, it is generally believed

P43-10
11

that silica on the powder surface forms a protective layer and

provides more stability during dewaxing in an oxidizing atmosphere; while silicon carbide (in the presence of the excess

carbon necessary for densif ication) cannot tolerate an oxidizing

environment. Therefore, various methods of dewaxing should be

considered:

• Evaporation and oxidation at elevated temperature

• Pyrolysis in an atmosphere other than air

• Vacuum extraction at elevated temperature

• Solvent extraction.

The dewaxing rate is very dependent on the heating rate, temperature

range, partial pressure of gaseous species, part size, and total quantity in

an oven load. It is also imperative to devise a technique to reduce the total

dewaxing time. Finally, successful dewaxing of thick sections is a limiting

factor in ceramic injection molding.

Sintering and Product Evaluation

The ability to form complex shapes with silicon nitride in mass- production quantities at acceptable costs depends on finding an alternative to

the current hot pressing approach--one such approach is pressureless

sintering. However, silicon nitride, like many other nonoxide ceramics, is

difficult to sinter by itself because of its covalent bonding. In the early

seventies. Jack (1972) suggested that sialon-type materials (silicon aluminum

oxynitride —it is often loosely applied to any composition containing the

elements Si-M-O-N, M being the metal atom) are sinterable without applying

pressure. Prochasza et al (1978) have successfully densified SiC using only

small quantities of sintering aids. Deeley et al (1961) have shown MgO as one

of the most effective additives for hot pressing silicon nitride. Terwilliger

and Lange (1975) have shown that silicon nitride can be sintered in nitrogen

at atmospheric pressure to a density of 90 percent theoretical using 5 percent

MgO as sintering aid. However, they noted a considerable weight loss which

resulted fom the dissociation of silicon nitride. Mitomo (1976) and Priest

et al (1977) have been able to density silicon nitride in the presence of a

P43-11
12

sintering aid to 95 percent and reduce the dissociation using 1.0 to 2.0 MPa

of nitrogen overpressure. Giachello et al (1979) have densified silicon

nitride with MgO as sintering aid at atmospheric pressure and control both the

silicon nitride dissociation and MgO volatilization by embedding the parts in

a loose powder mixture. Meanwhile, the sintering mechanism involving MgO has

been studied extensively; its shortcomings are pointed out and alternative

sintering aids are proposed (Brook et al, 1977). Silicon nitride can be

sintered to densities greater than 90 percent theoretical using oxide or nonoxide additions and a nitrogen pressure from 0.1 to 8.0 MPa. Using an amorphous powder (SN 402) and 4 weight percent Y 2 O3 as sintering aid,

Rowcliffe and Jorgensen (1977) achieved 90-99 percent theoretical density in

(2 atm) N?. Priest et al (1977) achieved 90-95 percent theoretical density

using SN 502 powder and 20 weight percent CeO in 2.0 MPa N 2

. Masaki and

Kamigaito (1976) could density a silicon nitride plus 10 weight percent

MgAl 204 mixture to about 96 percent theoretical in 0.1 MPa N2

. Stewart,

Wills, and Wimmer (1977) densified a mixture of S

i

3

N4 -A

1

2 O3 -A 1 N to 98 percent

theoretical in 0.1 MPa N 2

. Greskovitch et al (1979) densified Si 3 N4 with

Be 3 N 2 plus Mg 2 Si additives or Be2 SiN 2 plus Mg 3

N 2 additives. For present

advanced gas-turbine programs, silicon nitride with the Y 2 O3 or Y 2 O3 -AI 2 O3

additive system is actively being developed (Gazza and Katz, 1980).

While a wide variety of furnaces are used for sintering studies, a graphite induction-heated furnace is most common. The sintering temperature

ranges from 1300 to 2000 C. Parts are usually placed in BN crucibles to

protect against carbon in the furnace atmosphere.

The extent of densif ication is usually determined from the density

and shrinkage data. The weight loss records any dissociation of Si 3

N4 at the

sintering temperature. Typical results are presented in Table 3.

We have used density and room temperature flexure strength as the

typical end properties to demonstrate the influence of process variables on the properties of sintered silicon nitride. Table 4 compares the density and

strength data of three injection molded silicon nitride parts with similar

composition and sintering temperature (1750 C/2 hr). In all three cases, the

density and strength are within a comparable range; also, this strength is

equivalent to the value reported by Galasso and Veltri (1981) for cold

P43-12
13

isostatical ly pressed material. Yet, better control on injection molding

process parameters would be necessary to achieve higher strength as obtained

with hot pressed, and machined silicon nitride of similar composition.

A fracture surface (Figure 9) indicates a large pore on the edge of

. > the specimen is the fracture origin. Such pores are associated in injection

molding with nonuniform dispersion of binder ingredients that are subsequently

removed in the dewaxing process.

SUMMARY

In this paper we have attempted to demonstrate that injection

molding of ceramic parts is a viable production process. The work described

indicates that while the process is viable, the basic knowledge on this

technology still leaves much to be desired. The five major steps in injection

molding ceramic parts are covered, using silicon nitride powder as the

material of interest. While ultimate parts characteristics and economics are

not achieved, at least many of the ramifications in each of the steps are discussed, and suggestions are promulgated toward achieving at least acceptable results in each step. Also touched upon are some testing methods, and

basic theory indicating the interrelationships of the steps and their processes. Molding machines and mold designs and their effect on ceramic parts

are touched on. Binder removal methods are indicated and sintering is touched

on as the final step toward useable ceramic injection molded parts.

Currently achievable density and strength data on silicon nitride

shapes that are presented indicate that the ceramic molding process is

technologically and economically feasible.

P43-13
14

REFERENCES

Brook, R. S., Carruthers, T. G., Bowen, L. J., and Weston, R. J., 1977. "Mass

Transport in the Hot Pressing of Alpha Si 3

N4

"

, Nitrogen Ceramics, edited by

F. L. Riley, p 383.

Deeley, G. G., Herbert, J. M., and Moore, N. C., 1961. Powder Met., 8 , 145.

Galasso, S. F., and Veltri, D. R., 1981. "Sintering of S

i

3

N4 Under Nitrogen

Pressure", AMMRC TR81-28, September 1979 - September 1980.

Gazza, G. E., and Katz, R. N., 1980. "Development of Sinterable Si 3

N4

",

DOE/AMMRC 1AG EC-77-A-01-1017.

Giachello, A., Martinengo, P. C., and Tommasini, G., 1979. "Sintering of S

i

3

N4

in a Powder Bed", J. Mater. Sci. 14, 2825-2850.

Gilmore, A. D., and Spencer, R. S., 1950. Plastics, 143-228.

Greskovitch, C., Prochaza, S., and Resolowski, J. M., 1979, "Basic Research on Technology Development for Sintered Ceramics", Technical Report, AFML-TR-76179,

Contract F-33615-76-6-5033, General Electric.

Jack, K. M., and Wilson, W. I., 1972. "Ceramics Based on the Si-Al-O-N and

Related Systems", Nature (Phys. Sci.) 238 (80) 28-29.

Mann, L. D., 1978. "Injection Molding of Sinterable Silicon Base, Non-oxide

Ceramics", Technical Report AFML-TR-78-200, September 1977 - October 1978.

Masaki, H., and Kamigaito, 0., 1967. "Pressureless Sintering of Si 3 N4 with

Additions of MgO, AI 2 O 3 , and/or Spinel, YOGYO-Kyokai-Shi, 84 (10), 508-12.

Metzner, B. A., 1959, "Flow Behavior of Thermoplastics", edited by Bernhardt, C. E., p 19.

P43-14
15

Mitomo, M., 1976, J. Mater. Sci., 11, 1103.

Priest, H. F., Priest, G. L., and Gazza, G. E., 1977, "Sintering of Si 3

N4 Under

High Nitrogen Pressure", J. Am. Cer. Soc., 60 (1-2), 81.

Prochaza, S., 1978. "The Role of Boron and Carbon in the Sintering of Silicon

Carbide", Special Ceramics, 6 , edited by P. Popper, p 171.

Qtiackenbush, L. C., French, K., and Neil, T. J., 1982. "Fabrication of

Sinterable Silicon Nitride by Injection Molding", Ceramic Engineering and

Science Proceedings, 20-28, Am. Ceram. Soc.

Rowcliff, D. J., and Jorgensen, P. J., 1977. "Proceedings of the Workshop on Ceramics for Advanced Heat Engines", Contract 770110UC95a-ERDA-Rep.

Stewart, R. W., Wills, R. R., and Wimmer, J. M., 1977. "Fabrication of Reaction

Sintered Silicon", J. Am. Cer. Soc. 60 (1-2) 65-67.

Terwilliger, G. R., and Lange, F. F., 1975. J. Mater. Sci., 11, 1169.

P43-15
Table 1. Characterization of SN 502 Si-jN^ Powder

Composition, percent

Amorphous S^N,^ 40. Alpha ^' 3^4 57.6

Beta S' 3 n 4

2.4 °2 2.70

Cl2

0.20

Si Trace Mean Particle Size Submicron

Surface Area, m2/g 4.389 Apparent Density, percent of theoretical 9.4 Tap Density, percent of theoretical 15.0 Compaction Ratio, percent 4.7

Table 2. Milling Conditions and Results for SN 502

Porcelain Mill

Milling Parameters

1890 cc

Si^N^, Y

2O3

, AI 2O3 Powder 250 g

AI

2O3 Cylinders, Coors High Purity 1.27 x 1.27 cm 2100 g Loading Ratio, by weight 5.4:1

Milling Aid, weight percent 1

Mill Speed 100 rpm Compact Density

Milling Pressing

Time, Pressure Milling Aids, Percent Theoretical Density

hr MN/m2 Stearic Oleic Naphthenic Triethanalamine

8 172 45.1 43.0 46.3 40.4

16

" 50.8 54.8 51.5 51.0

24

" 51.1 55.4 47.4 52.7

48 " 56.7 54.4 48.2 56.8

P43-1

C
Table 3. Sintering Results on SN 502 Si 3N 4 Vanes, Cups, and Test Bars

Sample

Density

After Burnout,

Percent of Theoretical

Density

After

Sintering, Percent of Theoretical

Shrinkage

After Burnout, percent

Shrinkage

After

Sintering,

percent

Weight

Loss During

Sintering,

percent

Vanes:

1 — 97.8 ND* 18.8 1.20 2 — 98.1 ND* 19.1 1.07

3 , , r-r., » 98.8 ND* 18.8 1.32 4 — 98.3 ND* 19.3 0.96 5 — 97.6 ND* 18.6 1.01 Cups:

1 — 97.5 ND* 18.5 0.91 2 — 98.4 ND* 18.4 1.11 3 98.2 ND* 18.2 1.18 4 — 98.6 ND* 18.6 1.14

5 — 98.5 ND* 18.5 —

Test Bars:

1 55.2 98.4 0.71 19.5 0.77

2 55.6 97.5 0.74 18.6 0.75

3 54.8 98.8 0.60 18.8 0.84

4 55.5 97.8 0.56 19.2 0.95

5 55.8 98.0 0.64 19.0 0.88

*ND = Not determined.

Table 4. Density and Room Temperature Flexure Strength

of Sintered Silicon Nitride

Source Composition

Fabrication MethodDensity,

g/cc

RT Flexure

Strength,*

MPa

Galasso and Veltri (1981) Si 3 N 4-15Y203 -3Al 203 Cold Isopressed 3.141 432 (3-Point) Mann (1978) Si 3N 4-8Y 203 -4AI 203 Injection Molded 3.09 402 (4-Point) BCL (unpublished) S^N^Y2O3

- 4AI 2O3 Injection Molded 3.04 380 (4-Point) Quackenbush et al (1982) si 3 N 4- 6Y 2°3'2Al 2°3 Injection Molded 3.13 385.8 (4-Point)

Unmachined.

P43-17
Figure 1. Typical Injection Molded Silicon Nitride Components

a. Rotor Ring b. Stator c. Supercharger Rotor

d. Combustion Prechamber Insert

Figure 2. SEM of As-Received GTE SN 502 Silicon Nitride Powder

Figure 3. SEM of SN 502 Silicon Nitride Dry Ball Milled With 1 Weight Percent

Stearjc Acid for 24 Flours

Figure 4. Types of Flow Behavior

Figure 5. Stages in Mold Filling with a Silicon Nitride Mix

Figure 6. Viscosity Versus Shear Rate for Various Mixes

6a. Polyethylene with Different S^N4 Loadings

6b. Effects of Modifying Agents on Si 3N4 Loading and Viscosity

Figure 7. Injection Molding Machines

7a. Screw Type Machine

7b. Plunger Type Machine

Figure 8. Molding Cycle

A, Dead Time; B, Mold-Filling Period;

C, Pressure Build-Up Phase;

D, Packing Period;

E, Discharge;

F, Cooling With Gate Frozen

Figure 9. SEM on Fracture Surface of Injection Molded Silicon Nitride

P43-18
COMPONENTS OF CURRENT INTEREST

INJECTION MOLDED SILICON NITRIDE ROTOR RING

1 a.

SILICON NITRIDE STATOR

P43-19
COMPONENTS OF CURRENT INTEREST

INJECTION MOLDED SILICON NITRIDE

SUPERCHARGER ROTOR

COMBUSTION PRECHAMBER INSERT

P43-20
CERAMIC POWDER

CHARACTERISTICS

l | 5^.m

AS RECEIVED GTE SN502 Si3N4 POWDER

Figure 2.

Figure 3.

P43-21
Apparent

Viscosity

APPARENT VISCOSITY AS A

FUNCTION OF SHEAR RATE FOR

VISCOUS FLUIDS

Figure 4.

P43-22
EFFECT OF VISCOSITY ON THE

MOLD FILLING CHARACTERISTICS

Figure 5.

P43-23

J
6a. Viscosity of Polyethylene at Different Levels of Silicon Nitride Loading

6b. Effects of Modifying Agents on Silicon Nitride Loading and Viscosity

Figure 6. Viscosity Versus Shear Rate for Various Mixes

P43-24
QC

UJ

a.

a.o

Figure 7a.

P43-25

HEATERS
Figure 7b.

P43-2C
Figure 8

P43-27
SEM PICTURE OF INJECTION MOLDED AND SINTERED

SILICON NITRIDE ON THE FRACTURE SURFACE

Figure 9

P43-28
DEVELOPING AN INFORMATION STOCKPILE

TO AID IN SUBSTITUTION PREPAREDNESS

Robert T. Nash

Vanderbilt University

Introduction

During the workshop the attendees had the opportunity to provide their

written views on the acquisition, organization and dissemination of

information describing alternative technologies which could be used if

shortages of critical materials occurred. Thoughtful comments were obtained

which in combination provide a coherent picture of the issue of stockpiling

information.

The following five elements of the information stockpile issue were

identified in the responses from the workshop participants which are:

° Types of Information Explaining Alternative Technologies

° Uses and Sources of Information Explaining Alternative Technologies

° Channels for Disseminating Information

° Influential Organizations

° Qualification of a Material

Each of these five elements will be discussed separately.

Types of Information

Explaining Alternative Technologies

There was a consensus that the following types of information should be

included in a substitution preparedness stockpile.

Case Histories by Application

Substitution Plans by Companies

Structure - Property Relations

Substitution Tables

Invariably, changes in processes are associated with any change in the use of

materials. Consequently, an alternative technology consists of both a change

P44-1
in material and the associated processes which are employed in the production

and subsequent fabrication of the material into a finished part.

In addition to technical and economic information the participants felt

that administrative information would be necessary. In particular, government

agencies should keep industrial organizations informed about plans which would

be employed during an emergency if a shortage of materials occurred.

It was recognized that much useful information is proprietary and hence

not likely to appear in any publication. Further, information concerning

processes is usually not amenable to publication in any but the most

rudimentary form. Hence, the workshop participants thought that it would be

useful to develop lists of people that are expert on alternative technologies

who could be contacted during an emergency.

Uses and Sources of Information

Explaining Technological Alternatives

An important distinction exists between changes in materials for which

qualification is required and those for which it is not. At one extreme a

change can be made in a tool steel through a manufacturer ' s decision. At the

other, materials used in nuclear power systems and aircraft can only be

changed after lengthy qualification procedures.

Only those technological alternatives which have undergone at least

limited application in an industrial organization would be used by industry

during an emergency. Therefore, a consensus supported the position that only

a technology which has actually been employed in serial production, or at a

minimum in prototype testing, should be included in the information stockpile.

Any changes in the use of critical materials during an emergency must be

simple, or they will not be implemented due to lack of time.

P44-2
Technological alternatives which are used by industry often originate in

government laboratories and universities. However, this work is only

technological information in process until applied to the manufacture of

actual materials and finished parts by industry. Laboratory results, no

matter how promising, would be of little or no value in an emergency.

Existing channels of information transfer should be employed to the

greatest extent possible in disseminating information on substitution

preparedness. The regular sources include:

° Journals

0 Conference Proceedings

0 Handbooks

° Government Publications

It was also the view of the workshop participants that these regular

sources of information should be supplemented. The principal means would be

to publish abstracts of papers already appearing in the open literature.

These abstracts could be made available through existing electronic

information retrieval systems as soon as they are prepared. In addition, hard

copy summaries of the abstracts could be published annually. When a suitable

body of information had been accumulated on a particular topic, publication of

a monograph or handbook would become appropriate.

A number of public and quasi public organizations were mentioned which

would play a role in the organization of a materials substitution stockpile.

These include:

Channels for Disseminating Information

Influential Organizations

NMAB

ASTM

MPC

AISI

Department of Commerce

Department of Interior

(Bureau of Mines)

Department of Defense

P44-3

J
To guide the development of an information stockpile, an advisory group could

be organized which would include representatives from industry and each of the

influential organizations listed above. This group could provide advice on

the most effective means of organizing a stockpile of information pertinent to

substitution preparedness.

Qualification of a Material

A crucial element in determining what information should be included in

an information stockpile is the role of qualification procedures. These

procedures help determine the length of time required to make any type of

change in the use of a material. Therefore, the role of qualification in

substitution preparedness should be evaluated carefully. The workshop

participants felt that an alternative technology should only be included in

the information stockpile after qualification of the material.

Findings Concerning an Information Stockpile

The following findings were obtained during the workshop concerning the

organization of an information stockpile.

1) Since qualification procedures greatly influence the period of time

required to change the use of a material, only that information

concerning materials which have been qualified should be included in the

stockpile where qualification is necessary.

2) Only that information concerning technologies which have been carried

through prototype testing should be included in the stockpile where

qualification is not necessary.

3) Existing channels for disseminating primary information should be

employed to the greatest extent possible. The existing channels should

P44-4
be augmented with Substitution Preparedness Abstracts which could be used

to locate primary information.

4) Those organizations which can influence the adoption of new type' ^f

materials should be instrumental in the development of the informatio..

stockpile.

5) Much information is never published. Hence, the names of people should

be made available who would provide advice on alternative materials and

processes for use during an emergency.

6) Information should only be included in the stockpile if the substitute

technology can be implemented within a reasonable period of time. If too

much time is required, the substitute technology will not be used in an

emergency.

P44-5

NBS-114A (REV. 2-8C) U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

1. PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No. 3. Publication Date

SHEET (See instructions) NBSIR 83-2679-2 July 1983

4. TITLE AND SUBTITLE

Technical Aspects of Critical Materials Use by the Steel Industry

Volume II B: Proceedings of a Public Workshop; "Trends in Critical

Materials Requirements for Steels of the Future; Conservation and

Substitution Technology for Chromium".

5. AUTHOR(S)

Edited by R. Mehrabian

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 COMPLETE ADDRESS (Street. City. State, ZIP)

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

U.S. Department of the Interior, Bureau of Mines

U.S. Department of Defense, Army Research Office

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 information. If document includes a significant

bi bl iography or literature survey, mention it here) This volume presents papers given at a public workshop sponsored

jointly by the National Bureau of Standards, Bureau of Mines, Army

Research Office, and Vanderbilt University. The workship "Critical

Materials Requirements for Steels of the Future; Conservation and

Substitution for Chromium," was held at Vanderbilt University, October

4-7, 1982; it featured 50 presentations by technical authorities from

steel producing and using industries.

Volume I of this publication draws extensively on these proceedings and

reviews technical opportunities for research in process improvement

and in alternative materiald development that would reduce U.S.

dependency in critical materials. The advanced technologies reviewed in

Volume I in addition to their implications for critical materials

conservation represent trends leading to better quality, lower cost

steel products, and therefore they may contribute positively to the

industry's more immediate concern for improved markets.

12.

KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons)

Advance coatings; alternative materials; chromium; processing; steel;

strategic materials

13.

AVAILABILITY

| | Unlimited

[XI For Official Distribution. Do Not Release to NTIS

| Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402.

Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

415

15. Price

USCOMM-DC 6043-P80