NATIONAL INSTITUTE OF STANDARDS &

TECHNOLOGY

Research Information Center

Gaithersburg, MD 20899

NISTIR 89-3931

TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY APPLICATIONS

AT LOW TEMPERATURES - XII

j i 1 TEMPERATURE

Ic PrewrgJ.to

Department of Energy National Institute of Standards

Office of Fusion Energy and Technology

Washington, D.C. 20545 Boulder, Colorado 80303

NISTIR 89-3931

TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY APPLICATIONS

AT LOW TEMPERATURES - XII

Edited bv

R.P. Reed and R.L. Tobler

National Institute of Standards and Technolgy

Boulder, Colorado 80303

November 1989

Sponsored bv

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

^ flHIUNITED STATES

DEPARTMENT OF COMMERCE, Robert A. Mosbacher, Secretary

I I NATIONAL INSTITUTE OF STANDARDS

VHHP I AND TECHNOLGY, Raymond G. Kammer, Acting Director

CONTENTS

SUMMARY v

ORGANIZATIONAL CONTACTS vi

PRORGAM DESCRIPTION vii

STRUCTURAL ALLOYS

Structural Alloys Program 3

Austenitic Stainless Steels with Emphasis on Strength

at Low Temperatures 9

R.P. Reed

Nitrogen in Austenitic Stainless Steels 45

R.P. Reed

Effect of Nitrogen and Carbon on FCC-HCP Stability

in Austenitic Stainless Steels 65

R.P. Reed and M.W. Austin

Low Temperature Properties of High -Manganese Austenitic Steels ... 71

R.P. Reed, P.T. Purtscher, and L.A. Delgado

Discontinuous Yielding in Austenitic Stainless Steels

at Low Temperature 9 3

R.P. Reed and N.J. Simon

Influence of Interstitial Content on Fracture Toughness 107

P.T. Purtscher and R.P. Reed

Metallographic Study of the Crack- tip Region From Fracture

Mechanics Specimens of Austenitic and Ferritic Steels 119

P.T. Purtscher, R.P. Reed, and D.K. Matlock

Tensile, Fracture, and Fatigue Properties of Notched Aluminum

Alloy Sheets at Liquid Nitrogen Temperature 131

R. L. Tobler, R.P. Walsh, R.P. Reed, J.K. Han, and L. Ma

Elastic Constants of Four Fe-Cr-Ni-Mn Alloys 141

S. A. Kim, H.M. Ledbetter, and Y. Li

Nickel Effect on Elastic Constants of Fe-Cr-Ni Alloys 153

M. Lei and H. Ledbetter

Nitrogen Effect on Elastic Constants of F.C.C.

Fe- 18Cr- 19Ni Alloys 159

S. Lin and H. Ledbetter

iii
Low -Temperature Magnetic -Elastic Anomalies in F.C.C.

Fe-Cr-Ni Alloys 169

H. Ledbetter

Low-Temperature Thermal Expansion of Fe-Cr-Ni Alloys

with Various C+N Contents 177

M. W. Austin and H.M. Ledbetter

WELDING

Welding Program 183

The Fracture Toughness of 25Cr-22Ni-4Mn-2Mo Stainless Steel

Welds at 4 K 187

C.N. McCowan and T.A. Siewert

The Effect of Inclusions on the Fracture Toughness

of 316L Stainless Steel Welds at 4 K 199

C.N. McCowan and T.A. Siewert

TEST STANDARDS

United States — Japan Development of Test Standards Program . . . 217

U.S. — Japan Cooperative Program for the Development of 4-K

Test Standards 219

R.L. Tobler

Memorandum: Trip Report on ASTM D30 Symposium and Technical

Committee Meetings Related to the Development of Standards

for Mechanical Property Tests of Composites 223

R.L. Tobler

Charpy Tests Near Absolute Zero 227

R.L. Tobler and R.P. Reed

Warm Precracking at 295 K and Its Effects on the 4-K Toughness

of Two Austenitic Stainless Steels 239

R.L. Tobler and M. Shimada

Proposed Standard Method for Tensile Testing of Structural Alloys

at Liquid Helium Temperature 259

R.L. Tobler and R.P. Reed

TECHNOLOGY TRANSFER

Technology Transfer Program 279

Eleventh Cryogenic Structural Materials Workshop 281

N. J. Simon

IV
SUMMARY

This report contains the results of a research program to determine the

properties of materials that may be used in cryogenic structures for the

superconducting magnets of fusion energy power plants and prototypes. Its

purpose is to facilitate their design and development. The program was

developed jointly by the staffs of the National Institute of Standards and

Technology and the Office of Fusion Energy of the Department of Energy; it

is managed by NIST and sponsored by DOE. Research is conducted at NIST and

at other laboratories through subcontracts with NIST.

Research results for 1988 are presented in technical papers under four

headings that reflect the main program areas: Structural Alloys, Welding,

Technology Transfer, and United States-Japan Development of Test Standards.

Objectives and research highlights are summarized in the introduction to

each program area.

NOTE : Certain commercial equipment, instruments, or materials are

identified in this report to specify the experimental procedure adequately.

In a few instances, company names are used to identify the source of

specific research. In no case does such identification imply recommendation

or endorsement by the National Institute of Standards and Technology, nor

does it imply that the material or equipment identified is necessarily the

best available for the purpose.

v
ORGANIZATIONAL CONTACTS

Specific technical questions may be directed to the following people, who

contributed to major aspects of the program during the fiscal year of 1989.

Department of Energy, Office of Fusion Energy, Washington, D.C. 20545

Program Monitor D. Beard (301) 233-4958

National Institute of Standards and Technology, Boulder, CO 80303

Program Manager R.P. Reed (303)

(FTS)

497-3870

320-3870

Structural Alloys M.W. Austin (303) 497-3504

H. Ledbetter (303) 497-3443

P.T. Purtscher (303) 497-5789

R.P. Reed (303) 497-3870

N.J. Simon (303) 497-3687

R.L. Tobler (303) 497-3421

R.P. Walsh (303) 497-3818

Welding C . N . McCowan (303) 497-3699

T . A. Siewert (303) 497-3523

Technology Transfer N.J. S imon (303) 497-3687

E.S. Drexler (303) 497-5350

United States -Japan P.T. Purtscher (303) 497-5789

Development of R.P. Reed (303) 497-5789

Test Standards R.L. Tobler (303) 497-3421

R.P. Walsh (303) 497-3818

vi
PROGRAM DESCRIPTION

The objective of the program is to assist in the design, construction, and

safe operation of magnetic fusion energy (MFE) systems that use cryogenic

components, especially superconducting magnets. The specific steps taken to

achieve this objective are: (1) evaluation of the materials research needs

specific to MFE devices; (2) research programs to acquire the necessary

data; and (3) effective materials technology transfer by rapid dissemination

of the data to potential users through personal contacts, publications, and

workshops

Efforts directed at the first objective began with the publication of the

"Survey of Low Temperature Materials for Magnetic Fusion Energy" in March

1977. A publication updating part of this survey, "Structural Alloys for

Superconducting Magnets in Fusion Energy Systems," was included in Volume IV

(1981) of this series. In Volume VI (1983), reviews of the properties of

austenitic stainless steels and of their elastic constants also contributed

to this objective. Through interactions with low- temperature design,

construction, and measurement programs, such as the Large Coil Project, CIT,

and ITER (International Thermonuclear Experimental Reactor), we learn of new

problems as they arise. This year's report contains results of continued

research in support of Compact Ignition Tokamak designs by the Princeton

Plasma Physics Laboratory and the Massachusetts Institute of Technology.

Research projects contributing to the second objective are described in the

technical papers. Again, research toward the establishment of test

standards was emphasized. The third objective is satisfied, in part, by

these annual reports and by the series of workshops on Cryogenic Structural

Materials. Since 1982, handbook pages presenting the available data for

specific materials have been distributed to members of the magnetic fusion

energy community who are involved with low- temperature materials.

Vll

STRUCTURAL ALLOYS

STRUCTURAL ALLOYS PROGRAM

LEADER: R.P. Reed

CONTRIBUTORS: M.W. Austin, L.A. Delgado, S.A. Kim, H. Ledbetter,

P.T. Purtscher, N.J. Simon, R.L. Tobler, R.P. Walsh

OBJECTIVES

o Development of strong, tough structural alloys for use in superconducting magnets

o Development of strong, fatigue -resistant structural alloys for use in

pulsed superconducting magnets

o Development of effective codes and standards for low- temperature

property measurements and structural design

RESEARCH HIGHLIGHTS

o Two new reviews have been written:

Austenitic Stainless Steels with Emphasis on Their Low-Temperature

Behavior. Austenite, ferrite, body-centered cubic and hexagonal close

packed martensite, stacking- fault energy, carbides, and sigma phase

were described. Recent international development of higher strength,

tougher austenitic stainless steels was summarized. Discussions of the

effect of martensite phase transformations on the stress -strain

characteristics, nitrogen strengthening, and strengthening theory based

on lattice-parameter and elastic-property data, all at low

temperatures, were included.

Low- Temperature Properties of High -Manganese Austenitic Steels.

Emphasis was placed on strength, toughness, elastic properties, thermal

expansion, and magnetic properties, those properties most critical for

materials used in cryogenic structures. Strength and toughness

parameters and mechanisms were discussed, including alloying,

martensitic transformations, twinning, and grain size. Information was

included from many recent studies, particularly Japanese studies, of

alloy development for superconducting-magnet applications at 4 K.

o Regression analysis for significant design parameters has been carried

out on a matrix of 99 recent NBS measurements at 4 K on a variety of

316, 316 LN, and 316 LHN alloys. This matrix includes alloying,

refining, and processing data on alloys obtained from suppliers in

Japan, Europe, and the United States. For structural use in 4 K superconducting magnet applications, alloy 316 LN offers higher strength

than 304 LN with equivalent toughness. Nitrogen additions permit the

attainment of high yield- strength levels with little added cost.

3
o The fracture process at low temperatures in tension and from fatigue

precracking has been better characterized. Cleavage -like facets along

slip planes in austenite have been identified. Ductile dimples in the

fracture of tensile specimens have been related to inclusion

parameters. A new model has been developed for better understanding of

void nucleation, which sometimes controls the fracture of austenitic

steels

.

4
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS

Alloy

[Designation]

Country

of

Supplier

Tensile

295-4 X

Elastic

295-4 X

Fracture

Toughness

295, 76, 4 X

Fatigue Crack

Growth Rate

295, 76, 4 X

Fatigue

S-N

4 X

Report #

Voluae

(pege)

AUSTENITIC STEELS

Fe-5Cr-25Mn-0.2C Japan VII VII (65)

Fe-5Cr-26Mn Japan VI VI (181)

F«-13Cr-19Mn USSR III VI III (79)

VI (41)

Fe-13Cr-22Mn-0.2N Japan VI VII (65)

Fe-16Cr-8Ni-8Mn U.S. VIII VIII VIII (151,181)

Fe-17Cr-18lfa-0.5H U.S. VII, VIII VII (65)

VIII VIII (181)

Fe-1 7Cr-9Ni-8Mn U.S. III III (91)

Fe-17Cr-13Ni-2Mo U.S. I. II, VI 1,11, II, VIII, I I (15,71)

[316] III. IV, VIII, IX II (79)

VI, VIII, IX,X,XI III (49,105,117)

IX,X,XI IV (147)

V (185)

VI (157)

VIII (181,209,251)

n (15,27,43)

X (77)

XI (71)

Fe-17Cr-13Ni-2Mo U.S. II, IX, XI II, IX, XI II II (79)

[316LN] IX (15,27,53)

XI (71,103)

Fe-1 8Cr-( 11-14 )Ni- U.S. X X X X (9,31,43,55)

(0-4)Mo

[316LN]

Fe-18Cr-3Ni-13Mn U.S. i, VI, III I, VIII, XI I, VI, VII I (93)

VIII, XI III (91)

VI (53)

VII (85)

VIII (167,219)

XI (89)

Fe-18Cr-8Ni- U.S. X X (67)

(0-6)Mn

Fe-19Cr-(6-15)Ni- USSR VIII, XI VIII, IX, XI VIII (123)

4Mn-(0.1-0.3)N IX (53)

XI (113)

Fe-19Cr-9Ni U.S. I.II, I.II, I.II, II, III, IV I (15,71,213)

[304] (cont'd III, III, III, II (79,149,175)

on next page) III (15,91,105,117)

‘MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-884 (I);

NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (17); NBSIR 82-1667 (V); NBSIR 83-1690 (VI);

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX); NBSIR 87-3067 (X); NBSIR 88-3082 (XI)

5
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STS. URAL ALLOYS, continued

Alloy

[Designation]

Country

of

Supplier

Tensile

295-4 X

Elastic

295-4 X

Fracture

Toughness

295, 76, 4 X

Fatigue Crack

Grovth Rate

295, 76, 4 X

Fatigue

S-N

4 X

Report #

Volume

(P«*«)

Fe-19Cr-9Ni U.S. IV, VI, IV, VIII IV, IX, IV IV (37,101,203,

[ 304 ] , cont '

d

VII, IX,

X

X 215,227)

V (71)

VI (73)

VII (13,47,75,15:

IX (27)

X (77)

Fe-19Cr-9Ni u.s. II, III, I, II, IV II. IV III, IV, I (213)

[304L] IV, VIII, VIII, IX VIII II (79,123,175)

IX, X IV (37,101,131,

215,302)

Till (181,229,251]!

3 (27,141)

X (9)

Fe-19Cr-9Ni- U.S. IV, V, V IV, V, VIII, IV (77)

(l-lO)Mn VIII IX V (15,59,189)

VIII (181)

IX (53)

Fe-19Cr-9Ni-N u.s. I, III, IV I. Ill, IV, I, III, IV I 93)

[304N] IV, IX IX III (15)

IV (37,101,203,

215)

IX (27)

Fe-19Cr-9Ni-N u.s. II, III, IV, VIII II, III, IV, II, III, IV, II (35,79)

[ 304LN] IV.V.VI VIII, IX VIII III (15)

VIII, IX Iv (37,101,203)

V (29)

V (113)

VIII (181,207,229)

IX (159)

Fe-19Cr-10Ni- u.s. III, VIII VIII, IX, III, VIII III III (15)

( 0.03-0. 29 )N- XI VIII (145,181)

(0.03-0.09)C IX (159)

XI (131)

Fe-20Cr-7Ni-8Mn u.s. VIII VIII VIII (151,181)

[216]

Fe-20Cr-8Ni-2Mn u.s. VIII VIII VIII (151,181)

[304HN]

Fe-20Cr-16Ni-6Mn USSR V,VI V V (29,213)

VI (113)

Fe-21Cr-6Ni-9Mn u.s. I, VIII III, VI I, VIII I I (15,71)

III (91)

VI (15”'

VIII (1? 31)

VIII (1 i;

IX (95)

X (9)

‘materials STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-884 (I);

NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-1690 (VI);

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX); NBSIR 87-3067 (X); NBSIR 88-3082 (XI)
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

Country Fracture Fatigue Crack Fatigue Report 4 Alloy of Tenalle Elaatlc Toughness Growth Rate S-N Volume

[Designation] Supplier 295-4 X 295-4 X 295, 76, 4 X 295, 76, 4 X 295, 4 X (p«8«)

Fe-21Cr-12Ni-5Mn U.S. I. II III I. II I, II I (93)

II (79)

III (91)

Fe-25Cr-21Ni U.S. v.vi. III, VI V,VIII III (105,117)

[310] VII, V (71.145)

VIII, VI (73)

IX, X VII (65)

Fe-19Ni-9Co

[200-300 grade*

]

U.S. IV IV (237)

Reviews of IV.V.VI, IV.V.VI, IV.V.VI, IV.V.VI, VII IV, IV (17,257),

auatenltlc steela XI XI XI VI V (171)

VI (11,127,157)

VII (103)

XI (3,47)

ALUMINUM ALLOYS

Al-2.7Cu-2.2Li0.12Zr

U.S. X X (163)

Al-6Cu-0.2Mn-0.1Fe

[5083-0]

U.S. II II II II II (19)

Al-0.6Fe-0.2CuO.lSi

[1100]

U.S. IX IX (107)

Al-4Mg U.S. IV II II (35)

[5083-0] IV (185)

COPPER ALLOYS

99.99Cu-Ag U.S. IX, XI IX (95,107)

[C10100, C10200] XI (147)

99 . 5Cu-Ag

[C10400]

U.S. X X X X (83,127,139)

99.75Cu-Ag U.S. X X X X (83,127,139.

[C15500] 151)

9 7 . 6Cu-Ag-0 . 38Be

[ Cl 7510

]

U.S. X X (83)

99.95(Cu+Fe)-18Cr13Mn-3Ni-0.37N

[laalnate]

U.S. XI XI (159)

SUPERCONDUCTORS

Nb-45Ti U.S. III III: (133)

* MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78 -884 (I);

NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (17); NBSIR 82-1667 (V); NBSIR 83-1690 (VI);

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX); NBSIR 87-3067 (X); NBSIR 88-3082 (XI)

7
.

-
AUSTENITIC STAINLESS STEELS

WITH EMPHASIS ON STRENGTH AT LOW TEMPERATURES

Richard P. Reed

National Institute of Standards and Technology

Boulder, Colorado

The structure and properties of austenitic stainless steels

are reviewed, with emphasis on their low- temperature behavior.

Austenite, ferrite, body-centered cubic and hexagonal closedpacked martensite, stacking- fault energy, carbides, and sigma

phase are described. Recent international development of higher

strength, tougher austenitic stainless steels is summarized. This

chapter discusses the effect of martensite phase transformations

on the stress—strain characteristics, as well as nitrogen

strengthening and strengthening theory based on lattice -parameter

and elastic-property data, all at low temperatures.

INTRODUCTION

Rustproof steels, with increased passivity, were discovered at the turn

of the century. These new steels, called stainless steels, are characterized by high Cr (> 10%*) and low C (< 0.3%). Research beginning in 1904

by Guillet (1904—1014, France), 1 followed by Portevin (1909—1912, France), 2 Giesen (1909, England), 3 and Monnartz (1911, Germany) 4 led to an understanding of the corrosion behavior of Fe—Cr—C alloys. During the same

period, Guillet (1906) 5 and Giesen (1909) 3 published studies of Fe—Cr—Ni

austenitic stainless steels. Clearly, the objective of these early studies

was to develop a rustproof or corrosion- resistant alloy.

Since then, an enormous amount of research has been directed toward

understanding and improving the properties of Fe—Cr—Ni alloys. From their

initial use as cutlery, the applications of these steels have broadened

considerably, along with their compositions. Only two criteria determine

whether steels are classified as stainless (or rustproof or pitless)

Fe base element and about HCr or more. This broad classification is

divided into six principal subclasses:

^Weight percent is used, unless otherwise specified.

9
1 . Ferritic steels contain 11 to 30Cr, have little or no Ni, and low C.

Ferritic structure is body-centered cubic (bcc) and ferromagnetic ; it

is usually obtained by slow cooling from the hot forming temperature.

2. Martensitic steels typically have high Cr and low C and Ni contents.

Martensitic structure is bcc; its greater hardness and strength are

obtained by quenching.

3. Austenitic steels typically contain 18Cr and 8Ni. The face-centered

cubic (fee) austenitic structure is stabilized by the Ni content.

4. High -manganese . austenitic steels typically contain 24Mn, instead of

8Ni, to stabilize the austenitic structure. Their development was

driven by the relative scarcity and high costs of Cr and Ni.

5. Duplex steels contain varying amounts of Ni (4-8%) to adjust the

relative amounts of ferrite and austenite and, thus, to achieve high

strength (from the ferrite) and adequate toughness (from the austenite)

6. Precipitation-hardened steels contain Ai , Nb , and Ti, which form

carbides during aging at high temperatures. The carbides precipitate

in either martensitic or austenitic structures, increasing the strength

at the higher temperatures

.

The ternary Fe—Cr-Ni and Fe-Cr—Mn diagrams in Fig. 1 show the chemical

boundaries of these elements for the six types of stainless steels. Often

these steels are strengthened by adding interstitial C or N or the solidsolution alloying element Mo.

The earliest book (1935) on stainless steels that has come to our

attention is a series of technical papers on production and fabrication,

properties, test techniques, and applications edited by Thum. 6 Then in 1949

Zapffe produced a simplified book7 that emphasized the three structures

(ferritic, martensitic, and austenitic) of stainless steels. This book also

contains the fascinating history of the development and industrial manufacture of these steels. Two volumes devoted to stainless steel applications,

production, corrosion resistance, properties, and alloying effects were

written by Monypenny in 1951 and 1954.®

9 In 1956, Keating wrote a useroriented book 10 on Cr-Ni austenitic steels that focused on properties and

fabrication techniques.

The American Society for Metals published Source Book on Stainless

Steels (1976) 11 that includes property, fabrication, corrosion, forging and

heat -treatment, phase -diagram, and other information central to the use of

stainless steels. To date, the exhaustive handbook assembled by Peckner and

Bernstein in 1977 13 represents the most comprehensive review of these

alloys; it documents stainless steel melting, fabrication, metallurgy,

corrosion resistance, properties, and applications. Lula recently (1986)

revised a general introductory book by Parr and Hanson, 18 in which they

discuss metallurgy, properties, corrosion, and fabrication of stainless

steels. In the same year, he also wrote a careful review of the metallurgy,

properties, and applications of high-Mn stainless steels.

10
90

(Top) Fe-Cr-Ni ternary alloys.

(Bottom) Fe-Cr-Mo ternary alloys.

Figure 1. Composition ranges of various types of stainless steels.

From the 1950s into the 1980s, several major international conferences

have been held on stainless steels in general or specific aspects of these

alloys. Those of which we are aware are:

1.

Symposium on o-phase formation and verification; 1950;

sponsored by the American Society for Testing and

Materials

15

2. Symposium on the residual element effects on properties;

1966; sponsored by the American Society for Testing and

Materials

16

3. Conference, Stainless Steels for the Fabrication and User;

1969; proceedings published by the Iron and Steel Institute,

London. 17 At this conference, some of the first reports of

dual -phase stainless steels were presented.

11
4. Conference on stainless steels in general, with emphasis on

new applications; 1977; sponsored by Climax Molybdenum

Company and held in London. 1 *

5. Symposium on stainless steel castings; 1980; sponsored by

ASTM. 19

6. Conference on stainless steels with emphasis on properties

and alloy development for cryogenic applications; 1982;

sponsored by the International Cryogenic Materials Conference

and held in Kobe, Japan. 20

7. Conference on stainless steels with emphasis on fabrication

techniques, selection criteria, new developments, and new

applications; 1985; proceedings published by the American

Society for Metals. 21

Thus, notable stainless steel development, applications, and associated

research have been significant for about fifty years. During this time,

stainless steels have been used in many applications with material requirements far exceeding the definitive rustproof property of the steals. They

are versatile because they have high toughness, variable strength, high

elastic moduli, excellent weldability, and other practical properties.

Their stainless nature, of course, ensures low maintenance costs in most

applications

.

Austenitic stainless steels are usually chosen for cryogenic applications by virtue of their low thermal conductivity, good weldability, and

excellent toughness. They are well suited for applications that require

unusual safety considerations, such as storage of some liquefied gases.

Currently, stronger, tougher austenitic steels are needed for use at very

low temperature, 4 K. Nitrogen- strengthened Fe-Cr-Ni and Fe-Cr-Mn steels

are being studied and developed, especially in Japan.

Subsequent sections review how alloying affects the general structural

characteristics and properties of stainless steels. The first section

discusses their solidification, precipitation, and transformation

structures; the second section presents the effects of alloying on their

cryogenic properties.

STRUCTURES

Stainless steels are complex. They solidify as either an austenitic

(fee) or a ferritic (bcc) structure. At intermediate temperatures, an

intermetallic compound ( o

- phase ) , carbide precipitates, or both may form.

During cooling or plastic deformation, two martensitic transformation

products with bcc and hexagonal close-packed (hep) structures may form.

At least three magnetic phases are possible: austenite is typically paramagnetic at high temperatures and antiferromagnetic at very low temperatures; ferrite is ferromagnetic.

12
Austenite

Austenite is naturally the dominant phase of austenite stainless steels

at ambient temperatures. Austenite is sometimes called 7 and is always

designated 7 in phase diagrams. Solid-solution alloying elements, such as

Cr, Ni, and Mn, assume normal lattice sites in the Fe-base, fee structure;

there is no convincing evidence of ordering of these elements in the austenitic steels. Smaller atoms, such as C, N, P, and S, are thought to be

located interstitially in fee octahedral lattice sites. For these elements

(especially N) , electronic bonding or short-range ordering may be present,

but so far the evidence is circumstantial. Austenite remains paramagnetic

on cooling until about 50 K; then it usually becomes antiferromagnetic. The

magnetic structure of alloys with high-Ni content may become mictomagnetic

or spin-glass at low temperatures. The N6el temperatures, or temperatures

of transition to mixed ferromagnetic states, are strongly dependent on

alloying.

FeyrLEfi

Ferrite may form during solidification or high- temperature treatment,

depending on alloy content. It is occasionally present in small amounts

(usually less than 10%) in alloys that are rapidly cooled, such as weldments. This bcc phase is conventionally called 6 -ferrite when it forms at

temperatures above the 7 - loop or range and a-ferrite when it forms at lower

temperatures. Most austenitic stainless steels have sufficient stabilizing

elements, such as Ni, Mn, C, and N, to lower the temperature of the 7 - loop

or range sufficiently to prevent the diffusion-controlled t-*o transformation. In wrought alloys, the retained 6 -ferrite is typically transformed to

austenite during forging or hot rolling. However, since the solubilities of

many elements differ between the two phases, the solidification structure

affects the location of these elements. This is especially significant for

elements, such as P and S, that are more soluble in ferrite; therefore, they

are less prone to precipitate along grain boundaries during solidification

when 6 -ferrite is present. For this reason, 6 -ferrite is thought to avert

hot cracking in austenitic stainless steel weldments.

The relative amounts of austenite and ferrite that form during solidification are critically dependent on alloying. The liquidus and solidus

surfaces of the Fe-Cu-Ni ternary have been presented by Speich22 and

Schurmann and Brauckman . 2

9

The eutectic trough extends from near the

Fe comer (76Fe—10Ni-14Cr) to a ternary eutectic (about 8Fe-43Ni-49Cr) . Primary austenite solidifies if the composition is on the Ni side of the

eutectic liquidus. Nearer the eutectic liquidus on the Ni side, limited

ferrite forms, mainly in the dendritic cell boundaries. On the Cr-rich side

of the liquidus line, primary ferrite solidifies . 24 Upon cooling, however,

ferrite transforms to austenite in compositions typical of austenitic stainless steels. The ternary isothermal sections at 1350*C and 1100°C are shown

in Fig. 2 . At lower temperatures, the 6+7 region increases with attendant

decreases in both 6 and 7 single regions.

13
Chromium

F“"r* 2

- isTM; u-"> -

The tendency to transform from 5-ferrite to ^ is i 1 u

^

m? t"t f* amm b

“ -*i« J x“TMS«y

». <!!»> ;; s csr y ~ ,op

; •• ->««

“J““- by snail sdsuisn, * . ® “““ “*

CHROMIUM, wt.%

Figure 3. Binary Fe-CrxNi S0 (a) and

depicting dependence on Cr

of (a) 30 and (b) 40 wt.%.

CHROMIUM, wt.%

e-CryNi 40 _y (b) phase diagrams

and Ni at constant Fe contents

14
lit 5

<c

lit I

UJ

Figure 4. Binary Fe-Cr and Fe-Ni phase diagrams.

The binary Fe—Cr and Fe—Ni phase diagrams are shown in Fig. 4. Whereas

adding Cr encourages ferrite, adding Ni stabilizes the austenite. Indeed,

in both binaries, there is an extensive solid-solution range (>85% alloy

contents) at temperatures above 800 °C (below 800° C, a -phase forms in Fe—Cr

alloys, and below 500°C, Fe—high-Ni alloys tend to order). Binary Fe-alloy

phase diagrams for Mo, Si, and Mn are reviewed by Novak . 25 Like Cr, Mo and

Si have extensive ferrite ranges; the 7 - loop extends to only about 2Mo and

2Si. Austenite is encouraged by alloying Fe with Mn; at 1000°C, the 7 -loop

extends to 70Mn.

For welding, the amount of 6 -ferrite present is normally estimated from

a Schaeffler 26 or DeLong27 diagram. The ferrite portion of this type of

diagram is shown in Fig. 5; the lines representing zero ferrite, or the

demarcation between pure austenite (above) or small amounts of ferrite

(below) , are plotted. Tie lines depicting constant ferrite amounts have

been developed empirically in terms of the Ni equivalent (austenite -forming

tendency) and Cr equivalent (ferrite -forming tendency). These equivalents

serve to identify the austenite-martensite and austenite—ferrite boundaries

at ambient temperatures, where many alloying elements are present.

The proposed Ni- and Cr-equivalent equations are summarized in Table 1.

Presumably, the Suutala formulations , 26 are the most reliable equations

because they are the most recent and were developed from the most data.

Manganese additions up to 2 to 4% have been found to promote austenite

formation; greater amounts tend to promote ferrite formation . 28 Nitrogen

additions result in greater austenite stability. Notice in Table 1 that the

more recent Hammar30 equivalents emphasize C more than N. Suutala concludes

that at low (normal) contents, Mn additions stabilize austenite, but at

higher contents (> 5%), Mn additions enhance ferrite formation. Thus Hull's

formulation , 28 which includes a negative Mn2 term, provides the best representation of Mn effects.

Ferrite is ferromagnetic at temperatures below about 600°C and paramagnetic at higher temperatures.

15
Figure 5. Ferrite portion of Schaeffler diagram depicting lines that

estimate the first solidification of ferrite during cooling to

room temperature after welding.

Table 1. Ferrite Solidification Equivalent Coefficients.

Cr Equivalent Coefficient

(per wt.%)

Investigator (year) Cr Mo Si Nb Ti Ai

DeLong et al. (1956) 1 1 1.5 0.5 — —

Hull (1973) 1 1.21 0.48 0.14 2.2 2.48

Hammar, Svensson (1979) 1 1.37 1.5 2 3 —

tfi Equivalent Coefficient

(per wt.%)

Investigator (year) Ni Ml C N Cu Co Mn2

Delong et al. (1956) 1 0.5 30 30 — — —

Hull (1973) 1 0.11 24.5 18.4 0.44 0.41 -0.0086

Harmar, Svensson (1979) 1 0.31 22.5 14.2 1 — —

16
Martensite

The use of austenitic stainless steels is complicated by the metastability of the austenitic structure of most alloys. The metastability

leads to martensitic transformation, which is a significant design consideration in applications requiring fracture control planning, close

dimensional tolerances, the absence of a ferromagnetic phase, and high

toughness of weldment, heat -affected zone, and base metal. Either upon

cooling, under applied elastic stresses, or during plastic deformation, the

austenitic phase of Fe-Cr-Ni alloys may transform to bcc (o') and hep («)

martensite phases. The alloy composition affects the chemical free-energy

difference between the two phases.

Kaufman31 derived free-energy relationships between 7 and a' for the

Fe-Cr-Ni ternary system. Reed32 reviewed the Fe-Cr-Ni ternary free-energy

relations between fee and bcc and suggested a relationship for the freeenergy differences between fee and hep structures based on the earlier

binary Fe—Cr- and Fe-Ni-alloy work of Kaufman . 31

Other alloying elements, such as C, N, Mn, Mo, and Si, contribute to

austenitic stability in austenitic stainless steels. The complexity of

the final structure has led to many studies to characterize the effects

of alloying on the stability of the austenitic structure with respect to

a' -martensite

Expressions relating the stability of the austenite and the temperatures of transformation during cooling (Tms ) or during deformation (Tmd)

have been developed empirically and are summarized in Table 2. Eichelman

and Hull , 33 working with austenitic alloys with compositions of 10 to 18Cr,

6 to 12Ni, 0.6 to 5Mn, 0.3 to 2.6Si, 0.004 to 0.129C, and 0.01 to 0.06N,

established that all these alloying elements stabilize the austenite and

thus lower Tms . Monkman et al ., 34 using a larger of number of specimens

with compositions of 5 to 13Ni, 11 to 19Cr, and 0.035 to 0.126(C + N)

,

produced a similar analysis and concluded that Tms was linearly dependent

on composition only to a first approximation. The dependence of Tms

on the C and N concentration seemed to be influenced by the Cr and Ni

concentration.

Hull 29 studied the effects of Ni, Cr, Mn, C, N, Si, and Co (Ai, Th,

V, and V sometimes were added) on a' -martensite formation during lowtemperature cooling-and-deformation experiments and on ferrite retention

after cooling to room temperature from the melt. Hull's study confirmed

that all the above elements suppress low- temperature a' formation during

either cooling or deformation. To predict the effect of elemental additions

on Tms , Hull assumed that the Ni contribution was the average of the Monkman

et al. and Eichelman and Hull results and then compared all the effects of

alloying additions with that of Ni. The formulation of Andrews 36 applies

to lower Cr and Ni concentrations , and thus higher Tms ; his results are

included in Table 2 to indicate the disparity between the high- and lowtemperature empirical results. In all studies, specimens of various

17
compositions were prepared and cooled, and their Tu values were measured.

Regression analyses were used to obtain the empirical dependence of Tms on

alloy concentration. Self* • recently studied a vide range of compositions

of veld deposits. For analyses, he assumed nonideal behavior, and vith

regression analysis, he derived a predictive equation containing interactive

Cr-Ni and (Cr + Ho)-C terms (Table 2).

From Tables 1 and 2, it is apparent that the roles of Cr, Ho, and

Si are reversed in the stabilization of the bcc structure (ferrite and

a' -martensite) betveen high and lov temperatures. Binary Fe-Cr free-energy

derivations* 1 * 32 do not suggest this trend. Perhaps the explanation for the

stabilizing effect at lover temperatures is the contribution of these elements to the increase in austenite flov strength. The effect of Ho on flov

strength is discussed later in this paper.

The analytic expressions of Schaeffler, Eichelmann and Hull, and Self

are plotted in terms of Ni and Cr equivalents in Fig. 6. The less conservative analysis of Self implies that at room temperature, considerably less

than 18Cr is required to achieve the complete austenite stability of an

8Ni- equivalent alloy. Also, the nonlinear nature of the Self equivalent

indicates the possibility of martensite control by means of alloy optimization rather than alloy "trade-offs," as offered by the linear trend lines.

Table 2. Temperature equivalent coefficients for calculation of Tms

Element

ivalent coeff: 1 rqu LCiisiL, tyVC.% i

Eichelmsn,

Hull

(1953)

Montanan,

Guff, Grant

(1957)

Andrews Hull

(1965) (1973)

Base . . 1578 1455 273 1755 794

Cr ... -41.7 -36.7 -12.1 -47 -14.3

Ni ... -61.1 -56.7 -17.7 -59 -17.3

Mo ... -33.3 — -30.4 -54 -28.9

Si ... -27.8 — — -37 -37.6

C . . -1670 -1460 — -2390 -350

N . . -1670 -1460 -423 -3720 —

Mo — — -7.5 -56 -29.6

Ocher — -180(Ti) -1.19 (CrNi)

+23.1 (Cr+Mo)C

18
NICKEL EQUIVALENT

For constant amounts of 1 Mn. 0.03 C, and 0.5 Si. Figure 6. Martensite portion of Schaeffler diagram depicting lines that

estimate the first formation of martensite during cooling to room

temperature after welding.

Experiments have been conducted to assess the influence of cold work

(either tensile or compressive) on Tm(j. Results are contained in Table 3.

The experimental definition of Tm(j varied in each study: Angel 37 used 50%

a' at 30% tensile elongation; Williams et al. 38 used 2.5% a' at 45% compressive elongation; and Hull 29 used the minimum detection level (probably about

1% a') at 50% compressive elongation. There are subtle distinctions among

the coefficients for Tms and Tm<j calculations. After cooling, the Cr and Ni

coefficients are nearly equal, Mn is higher, C and N are a factor of 50

higher; after deformation, Cr and Mn are decidedly less than Ni, and C and N

are about a factor of 10 higher. Careful research to delineate the dependence of composition on free energy, stacking- fault energy, and deformation

parameters is needed to understand these distinctions.

Table 3. Temperature Equivalents for Calculation of Tm<j

19
Upon cooling or during deformation, fee austenite may transform to two

martensitic products. The hep phase is associated with extended stacking

faults and forms as thin sheets on (111) austenite planes. In most austenitic stainless steels, the a ' product forms as laths. Breedis 39 reported a

lath a' morphology for compositions ranging from Fe-19Cr-llNi to Fe-lOCr16Ni; at lower Cr and higher Ni concentrations the a' morphology changed

to surface martensite and then to a plate-like structure (Fig. 7). From

Reed , 40 the a' lath-like structure is parallel to <110>7

, with (225)~,

( 1 1 2 )

-y , or both habit planes . 41 The laths are restricted within {111}^

bands, and usually three sets of habit planes form within the band.

There is a considerable amount of accommodation deformation in the

austenite after transformation during cooling, particularly within the { 111 } bands containing the a' laths. Either e-martensite or large amounts of

stacking faults are present.

The orientation relationships between the 7 -, e-, and a' -phases 41 are

( 111)7 || ( 0001 ), || ( 101 ) o

.

[ 110 ], || [ 1210 ], || [ 111 ],,.

( 1 )

These relationships are apparently retained regardless of the manner of e or

o' formation.

Transformation during plastic deformation initially results in

a' -martensite at intersections of active (111) deformation bands. The

a' assumes the shape of laths along <110> common to the two active ( 111 ) systems. Evidence of dislocation pileups at a' laths has been observed . 32

Also, more than one lath forms at each transformation site, and in contrast

to the a' -phase that forms on cooling, all laths have the same habit-plane

variant.

Figure 7. Compositional dependence of a' -martensite morphology in Fe-Cr-Ni

ternary alloys.

20
The internal defect structure of the o' -phase in Fe—Cr—Ni steels consists predominantly of dislocations. As Cr is replaced by Ni, Breedis 39

reports that the cellular, irregular distributions of dislocations that are

typical of a ' laths change to the planar, regular arrays that are typical

of a ' plates. Also, as Ni replaces Cr, the amount of e transformation decreases, the sharpness of the hep reflections decreases, and the fee tvin

reflections become more diffuse.

Many metallurgists have been concerned with the role of the € transformation: does € act as a precursor transformation (-)r»€-*a'), or is the « an

accommodation effect (tt*®' » T*0? Both effects have been observed. The

stainless steels with low stacking -fault energy tend to form e, and the a'

tends to form from c. Higher stacking -fault -energy alloys require transformation stresses for observable e formation.

Several characteristics of the austenite-bcc -martensite transformation

in stainless steels cause problems in many applications. The martensitic

bcc has a specific volume 1.7% larger than the parent austenite; therefore,

these steels expand during transformation. The martensite forms as individual crystals, and the shear stress and volume expansion associated with

transformation disrupt localized regions. For example, a' formation near or

at the surface results in localized surface upheavals. For service that

requires close tolerances, such as valves or bearings, these local surface

fluctuations are disastrous.

Another concern, mostly in the presence of magnetic fields or pulsed

currents, is that the bcc martensitic product is ferromagnetic, whereas the

parent austenite is paramagnetic. A simple rule is that each percent of bcc

martensite results in a permeability increase of 0.01. Therefore, if timedependent field changes corresponding to material changes of the order of

0.10 are significant, than alloy selection is important.

Stacking-Fault Energy

The stacking -fault energy is related to the free -energy difference

between the fee principal structure and the hep structure. Lower stackingfault energy leads to wider partial -dislocation separation (larger fault

ribbons), and hence, more planar dislocation slip structures and reduced

cross slip.

Schramm and Reed43 used x-ray peak-shift measurements , coupled with

previous stacking -fault energy measurements, to estimate the dependence of

the stacking- fault energy on chemical composition. This dependence should

correspond to the dependence of the transformation on composition,

because a stacking fault represents a local planar area of « . Rhodes and

Thompson43 suggested that the stacking -fault energy values of the Schramm

and Reed analysis were too large, considering additional weak-beam electronmicroscopy data. Weak-beam electron microscopy measurements by Bampton,

Jones, and Loretto 44 confirmed that the least-square analysis of Schramm

and Reed produced stacking- fault energies that were too large. From

measurements of individual nodes, they estimated a data spread of ± 25%.

Brofman and Ansell 46 added C content to the factors affecting stacking- fault

energy. The stacking -fault energy composition dependencies are contained in

Table 4.

21
Table 4. Stacking- fault energy at room temperature.

Energy Equivalent , mJ/(m2

• vt . %)

Investigator

(Year) Base Cr Ni Mn Si c Mo Other i

Dulleu, Shitting

(1964) 0 0.5 1.4 3.4 0.1 3.6 (Ti),

3.2 (Cu),

-0.55 (Co)

Schramm, Reed

(1975) -53 0.7 6.2 3.2 - — 9.3

Rhodes , Thompson

(1977) 1.2 0.6 1.4 17.7 -4.7 —

Brofman, Ansell

(1978) 16.7 0.9 2.1 26

Ledbetter and Austin46 recently examined with x-rays the effects of

additions of C and N (C + N 3 0.325 vt.%, N £ 0.21 vt.%) on the stackingfault energy of an Fe—8Cr—10N1 alloy. They reported an increase of about

10% in the stacking- fault energy per at.% C + N. Earlier, Stoltz and

VanderSande 47 determined with weak -be am microscopy that N in excess of

about 0.24 wt.% dramatically decreased the stacking- fault energy of an

Fe—19Cr—7Ni—8Mn-0.03C alloy. Fujikuma et al. 44 measured stacking- fault

probabilities for an Fe—18Cr—lONi—8Mn alloy. They reported a minimum at

about 0.15N; a slight increase at lover N contents, and a rapid increase

at higher N. Thus, all three sets of measurements appear to be consistent:

the stacking -fault energy increased slightly with N contents up to about

0.20 vt.%; it decreased substantially for higher N contents.

The solid- solution solubility of the mobile interstitials C and N in

austenitic alloys is very temperature dependent (Fig. 8). 49,60 At temperatures less than about 950*C, C begins to precipitate as a carbide in practically all austenitic steels, since the solubility limit is about 0.035.

Using electron microscopy, the solubility of C in type 316 steel (Fe-19Cr14Ni-2Mo) has been represented by the expression61 :

log [C (ppm)] - 7.771 - 6272/TOQ (1)

22
The solubility limits determined from Eq. (1) are considerably less than

those given in Fig. 8, which reflects the higher sensitivity of electron

microscopy for detection of carbides. Carbides precipitate in the range 500

to 1000 # C; below 500 # C the atom mobility is too low, and above 1000 # C the

solubility of C and N is sufficient. The most common carbide is M23 C 6 , where M represents a metal, normally Cr. However, Fe, Mo, Ti, and Nb also

may assume the position of the metallic element, and N and B may assume the

position of C in the M23 C 6 structure.

The solubility limits of both C and N are reduced by the presence of Ni

(see Fig. 8 for C) . On the basis of binary Fe—Ni data, 62 Fig. 9 shows that

this reduction is about 0.014N from the addition of 20Ni. The data of

Sakamoto et al., 63 obtained from observations of bubbling vacuum- induction

melts, are quite consistent and represent the practical limits of maximum

N contents (Fig. 9). These limits are particularly important to define the

N- strengthening limitations that are discussed under Strength . Consistent

with the binary alloy solubility limits, these data show that both Cr and Mn

increase the ingot capacity for N.

Carbides precipitate primarily at ferrite-austenite interfaces, followed by grain boundaries, noncoherent and coherent twin boundaries, and,

finally, within grains. This is depicted schematically in Fig. 10. 54 • 66

The temperature decreases and the time for nucleation increases as the

precipitation site changes from high- to low-energy boundaries. Increased

C content increases precipitation temperature and decreases time; increased

N content has the reverse effect, suppressing carbide formation. 12

- 1400

1300

1200

UJ

GC

<

GC

1100 W

CL2

UJ

1000 h

000

CARBON, wt.%

NITROGEN, wt.%

Figure 8. Solubility limits for carbon^ 49 ) and nitrogen

(

50 ) as functions of

temperature for various alloys.

23
X

UIs3

K

<

C

UI

Q.2

Ui

Figure 9. Solubility limits for nitrogen in Fe-Cr, Fe-Ni, and Fe-Mn binary

alloys at 1473 K^® 2 ^ and in Fe- 15Ni-Mn-Cr alloy liquids^ 63 ).

ALLOYING ELEMENT, wt.%

Figure 10. Time- temperature carbide precipitation kinetics, with

specific sites identified64 * 56 .

24
The M29 C6 carbide has a complex fee structure with 92 metallic atoms

and 24 smaller (C t N, or B) atoms and a lattice parameter of 10.61 to

10.64 A . 13 But there is only minor mismatch between the (111) M29 C 6 and

(111) austenite, estimated at 1.3% by Lewis and Hattersley . 66 Thus, it the

common plane of the Interface between the two structures has been identified

as (111 ). 64,56 Alloy additions to the M29 C 6 structure, such as N, tend to

decrease its lattice parameter and to increase the amount of interface

mismatch. This should retard the growth kinetics.

Carbide precipitates normally have a sheet- like morphology at lower

temperatures (480—730*C) and dendritic shapes at intermediate temperatures

(600-875*C); particles form at the higher temperatures (> 850 # C). 66 Higher

temperatures, as well as longer aging times, promote coarsened structures.

Carbide precipitation in austenitic steels has long been associated

with increased sensitivity to intergranular corrosion, termed sensitization.

For many applications, suppression of sensitization is achieved by rapid

cooling from solution- treatment temperature. During welding, however, the

heat-affected zone unavoidably enters the temperature range of carbide formation. To prevent severe sensitization during welding, two courses have

been followed: (1) low-carbon (^0.03) grades, designated by AISI with the

"L" nomenclature, have been provided and (2) alloys with Ti (AISI 321) and

Nb (AISI 347) have been developed to achieve more random TiC or NbC

precipitation within grains, rather than M23 C e precipitation at grain

boundaries. These alloys are often called stabilized grades of stainless

steel, since the C, taken out of solution, is unavailable for carbide

formation and precipitation at grain boundaries.

More thorough reviews of carbide formation in austenitic steels are

provided by Novak26 and Lai . 67

Sigma Phase

Notice the Fe-Cr part of the phase diagram in Fig. 4: at Cr contents

greater than about 15%, the o-phase may form when the alloy is exposed to

temperatures between 400 to 800*C. The o- phase extends into the Fe-Cr-Ni

ternary diagram, as indicated by Fig. 3. Alloys rich in Cr and lean in Ni

tend toward o-phase formation at lower temperatures. Maehara et al . 58

studied the effects of alloy additions on the amounts and aging times for

o-phase formation. They found that Ni decreases the amount of o-phase

formation, but increases the rate of formation in an Fe—25Cr—2 . 8Mo-base

alloy. Additions of Cr and Mo increase both the aging kinetics and amount

of o-phase in Fe-6.5Ni-2.8Mo and Fe-25Cr-6 . 5Ni-base alloys, respectively.

In Fe—Cr, the highest temperature of o-phase formation is 821 *C

(Fig. 4); Ni additions gradually increase this temperature (Fig. 3). The

constant-Fe vertical sections of the Fe—Cr—Ni ternary of Fig. 3 are consistent with the liquidus, solidus, 1100°C, and 650 # C isotherms of Speich 22 and

the 802 # C and 648 # C isotherms of Talbot and Furman . 59 Like carbides, the

o-phase forms preferentially at high-energy surfaces. Also, when ferrite is

present, there is a strong tendency for o-phase to form initially within the

ferrite or at the ferrite side of the 7—6 boundary.

25
The ?-phase has a tetragonal structure, 30 atoms per unit cell, an a0

lattice parameter range of -8.29 to 9.21 A, and cQ lattice parameter range

of -4.60 to 4.78 A . 13 Sigma-phase alloy constituents do not have fixed

stoichiometric ratios; they may range from B 4 A to BA4

. Alloy additions,

such as Cr, Mo, V, and Si, promote a -phase formation.

For a more extensive review, see Novak36 and Lai. 6T

PROPERTIES

Stress-Strain Characteristics

The AISI 300 series stainless steels used in cryogenic applications

range from metastable to stable austenites. Alloys such as 304, containing

18Cr and 8Ni, are metastable, whereas 310, containing 26Cr and 20Ni, is

stable with respect to martensitic transformation. The stress—strain

behavior and temperature dependence of the flow strength of these alloys

differ and depend on the austenite stability. In the less stable alloys,

both e- and a' -martensite form; in the slightly metastable alloys, such

as 316, only a ’ -martensite forms; and in the stable alloys, neither e- nor

a' -martensite forms during deformation to fracture at any temperature.

At low temperatures, the tensile stress-strain characteristics of

austenitic stainless steels depend on the stability of the austenitic

structure. Consider the engineering stress—strain curves of a stable austenitic steel, alloy 310 (Fig. 11a) and a metastable alloy 304 (Fig. lib) at

low temperatures. As the temperature decreases, the yield strength of alloy

310 increases significantly, and the work-hardening rates, reflecting only

dislocation interactions, remain relatively constant.

26
Figure 11. Low- temperature stress-strain curves for (top) a stable

austenitic alloy 310 (Fe-23Cr-20Ni) and (bottom) a metastable

austenitic alloy 304 (Fe- 19Cr-9Ni)

In a metastable austenitic, polycrystalline alloy at low temperatures,

three distinct stages are present in the stress-strain curves. The contours

of their stress—strain curves (Fig. 12) are remarkably similar to singlecrystal shear- stress—strain curves. Stage I represents the microstrain and

early macrostrain behavior. The formation of a' -martensite is not thought

to occur in this range; stacking -fault clusters, or c -martensite , or both

are most likely to complement dislocation interactions to reduce the rate of

work hardening.

Figure 12. Schematic depiction and characterization of three stages

associated with the stress -strain curves for polycrystalline

metastable austenitic stainless steels at low temperatures.

27
Stage II, the "easy glide" range, is associated with increasing

e-martensite formation and the formation of a' -martensite laths at crossslip intersections. Suzuki et al.*° proposed that such a' laths, with the ! long <110> direction representing the intersection of two active slip

systems, act as windows to assist cross slip. Another possible explanation

is that strain- induced e -martensite (whose formation increases prevalent

in this stage), contributes a larger strain component (or, alternatively,

exhibits lower work -hardening rates) . Stage II is most prevalent in the

temperature range 76 to 200 K. In this temperature region, the range of

strain at which the stress remains relatively constant is large, usually

from 0.04 to 0.10.

In stage III, the rate of work hardening increases and then becomes

constant over a relatively large amount of plastic deformation (20 to 40%)

.

The volume percentage of a' -martensite is linearly related to plastic

deformation in stage III. The nature of the transition in the roles of

a' -martensite formation — from possibly promoting easy glide in stage II t

being associated with a linear, high rate of work hardening in stage III —

is not clear. Perhaps in stage III essentially all active cross-slip sites

have been transformed to a' -martensite , and subsequent a' -martensite formation occurs within active glide bands. Such a' -martensite formation woul*

not be expected to promote stacking- fault glide. There is a hint from the

x-ray data of Reed and Guntner61 that the following sequence may occur:

A maximum value of the volume concentration of « -martensite occurs at

about the transition from stage II to stage III. There is apparently no

€ -martensite formation during stage III; the amount of « -martensite

decreases with strain, presumably transforming to a' -martensite.

Stages I and II decrease in significance in more stable austenites, bu

stage III is retained. For instance, AISI 316 exhibits neither stage I nor

stage II, but does begin to transform to a' -martensite at strains of about

0.02 below 190 K. No « -martensite was detected in 316LN alloys.

Alloying of austenitic stainless steels affects the stress and strain

parameters of the three stages of low- temperature deformation. The effects

of Mn alloying on these parameters at 4 K have been studied by Reed and

Tobler. 62 The yield strength, Oy (flow strength at 0.002 plastic strain) i

within stage I. As shown in Fig. 13, Oy is linearly dependent on Mn content, increasing with increasing Mn. The linear dependence is similar to

that of ojj, the stress at which stage II begins. The increase of Oy with

Mn is considerable, about 33 MPa per wt.% Mn. The addition of Mn to austenite solid solution does not alter the lattice parameter nor the shear

modulus significantly (see later discussion), therefore traditional solidsolution strengthening should not result from Mn addition. But, Mn

increases the stacking- fault energy and suppresses a'- and e-martensite

transformation; these phase transformations influence the stress and strain

parameters. Therefore, the increases in strength that occur when Mn is

added are due to Mn's contribution to austenite stability. Figure 13 shows

that an addition of 0.1N increases Oy about 290 MPa. Nitrogen has the same

strengthening effect in stable alloys, which implies that N has little

effect on hep or bcc stability within the composition range 0.1 to 0.2 wt.%

28
1.6

01 234 5 678

MANGANESE, wt. %

Figure 13. Dependence of stages I and II strength parameters on Mn and N

contents at 4 K.

The stress denoting transition to the region of high work-hardening

rates, ^xil» tracks very closely with ajj* The rate of work hardening (from

engineering stress—strain curves) of stage III decreases with increasing Mn

and N, as shown in Fig. 14. The addition of N slightly affects the dependence of Mn content; N contributes to the reduction of da/de about ten times

more strongly (on the basis of wt.% addition) than Mn. The dependence of

the work -hardening rate in stage III on the percent a' -martensite per unit

elongation is shown in Fig. 15. Clearly, there is a strong linear correlation, and the rate of work hardening can be regarded as a function of

austenite stability with respect to a' -martensite formation, as determined

by composition.

We have performed numerous correlations between the normalized amount

of strain- induced o' -martensite at final elongation at 4 K and the alloy

composition. Our best correlation for this normalized strain- induced martensite was obtained by using the Williams et al. 38 regression analysis

for Tm(j In Fig. 16 their expression is plotted versus percent o'/Percent

tensile elongation at 4 K for a series of modified 316LN alloys. Correlation is excellent for these alloys with a calculated Tm(j ranging from

120 to 240 K.

Therefore, the deformation characteristics of austenitic stainless

steels can be significantly controlled by alloying. Sufficient alloy

additions to stabilize the austenite result in typical solid- solutionstrengthened, fee stress-strain characteristics. The shape of the stressstrain curves of metastable austenites during strain- induced a'- and

c -martensite transformations is affected by alloying.

29
Figure 14. Dependence of stage 111 work-hardening rate on Mn and N

contents at 4 K.

Figure 15. Dependence of stage 111 work -hardening rate on normalized

amount of a' -martensite formed during uniform elongation at

4 K.

30
Figure 16. Correlation between normalized a' -martensite formation in

stage ill with alloying, as characterized by the Tmd

calculations of Williams et al. 37

Strength

In addition to austenite stability, the strength of austenitic

stainless steels is influenced by temperature, grain size, cold and hot

working, and interstitial and solid- solution alloying. Temperature and

grain-size effects have been discussed previously

63 ‘ 6S We focus here on

the effects of alloy strengthening on stable austenitic steels at low

temperatures

At very low temperatures, the variation of tensile yield strength (ay)

with N, C, Mo, Mn, and Ni additions to Fe—Cr—Ni austenitic steels has been

recently and extensively studied. In Japan a program to develop an alloy

with <jy of 1200 MPa and a toughness of 200 MPa-m^'^ has now been completed. 67—68 At NBS we have tested at low temperatures some of the new

Japanese alloys, the European 316LN alloys, and selected laboratory heats

from U.S. suppliers in which N, C, Mo, Mn, and Ni alloy additions were

systematically varied. A summary of these test results follows.

The interstitial additions substantially strengthen the alloys at low

temperatures

64 » 65 » 70 * 63 * In Fig. 17, the effect of C + N on Oy at 4 K for

an Fe—18Cr—lONi steel is presented. 72 The break in the linear dependencies

at about 0.15(C + N) is related to martensitic transformation; at lower

C + N contents, the strain- induced martensite reduces the lower flow

strength, as discussed in the Stress—Strain Characteristics section.

Regression analysis of 99 data points for alloy 316LN (and alloys with

similar compositions) at 4 K indicated that Oy increased 237 MPa per 0.1N; 77

an earlier analysis of 30 data points of 304LN-type alloys at 4 K found that

Oy increased 319 MPa per 0.1N. 70 Strengthening from C is about half that

31
of N at 4 K. 70 Thus, N additions strengthen considerably more than equivalent C additions at low temperatures. Alloying with N instead of C also

reduces the tendency toward sensitization, as discussed under Carbides . Furthermore, N is less expensive than C. For these reasons, N is now

usually selected to increase the low- temperature strength of austenitic

stainless steels. Many recent studies have measured the strengthening

contribution from N for both Fe-Cr-Ni and Fe-Cr-Mn base austenitic structures at 4 k.««»«6*7o- 76,77-83 Studies of Reed and Simon70 * 77 found a

smaller standard deviation from regression analyses using the power of one

to express the [N] dependence compared to 1/2 or 2/3 power dependence.

Fundamental theories for solid- solution strengthening84-86 predict dependence of the flow strength on concentration to the power 1/2 or 2/3 from

dislocations cutting through the stress field created by individual atoms

of this concentration. The linear dependence of Oy on C + N shown in

Fig. 17 also supports this assessment. This dependence of Oy on N alloying

is strongly dependent on temperature (Fig. 17). Reed and Simon70 found that

the N coefficient increased by a factor of 6 from 295 to 4 K, while the

coefficient for C increased by only a factor of 2. This strong temperature

dependence implies increased N bonding strength, or clustering, below room

temperature, since the temperature dependence of C (the larger atom), should

reflect the modulus and size effects (see later discussion)

.

The dependence of Oy on solid-solution alloying additions, such as Mo,

Mn, and Ni, is not as great. Molybdenum increases strength most effectively, about 50 MPa per wt.% at 4 k. 74 * 87-8 ® There is evidence that at

higher Ni contents the strengthening effects of Mo are reduced. 8 ® Mn, when

not affecting austenite stability, increases Oy by about 10 MPa per wt.%

at 4 k. S 3 » 73 * 76 * 8 ®*® 0 Nickel does not contribute directly to solid- solution

strengthening, but does affect strength through its strong contribution to

austenite stability. The temperature dependence of the solid- solution

strengthening additions has not been assessed.

200

*«

JC

180 b*

CO

CO

UJ

c

100 H

CO

o

-J

UJ

60 £

Figure 17. Dependence of yield strength on C and N contents in Fe-18CrlONi austenitic stainless steel at low temperatures.

CARBON +NITROGEN, wt.%

32
In Table 5, the strengthening contributions from substitutional and

interstitial alloying additions in fee elements 93 '* 7 and in austenitic

stainless steels 63 ' 91 are listed in terms of a fraction of the shear modulus. This refers to the increase of Oy from adding 100% (by atoms) of the

solute, found by extrapolating from available data that usually range from

1 for interstitials to 5 to 20 at.%.

Adding C to Ni 96 or C to austenite produces roughly the same increases

in strength at low temperatures. The addition of Mo to Ni or to Fe-Cr-Ni is

also approximately equivalent. Galligan and Goldman93 have achieved much

stronger strengthening effects from alloying by quenching- in Ag into

probable split- interstitial sites in Pb. In their work, strength increases

are linearly proportional to Ag content, not to a fractional power of it.

From the theory of solid- solution strengthening, 64 ' 66 the shear

strength, r, resulting from a random array of single atoms interacting with

mobile dislocations is

r= A[Cf

2/3] [f0V3], (2)

Table 5. Strengthening Contributions in fee Alloys.

Solvent Solute Temperature (K) Amount (Ao/100%)

Ag Ai 4 G/30

Sb 294 G/100

76 G/30

Ni Mo 295 G/45

C 4 C/7

295 G/30

Cu Ni 76 G/120

Au 76 G/100

Pb Ag* 4 2G

Sn 4 0

Fe-Cr-Ni N 4 G/2.5

(Austenite) 295 G/15

C 4 G/5.5

295 G/10

Mo 4 G/60

295 0

Mn 4 G/200

295 0

Ni 4 0

295 0

*Quenched from 300 K

33
where Cp is the concentration of e solute species, and A is considered to b<

a constant for this discussion. The maximum interaction force,

fS - n

2 +<**6*, (3)

where the modulus misfit,

6 1

G dc* (4)

In Eq. (4), G is the shear modulus, c is the concentration, and the size

misfit,

6 i &

a dc’

where a is the lattice parameter. The constant a is typically about 20. 84

Thus, the relative magnitudes of the strengthening contributions of various

solutes may be estimated by comparing the modulus and size misfit parameters, assuming that other possible strengthening effects (e.g., electrostatic, stacking- fault interactions) are small.

Ledbetter and Austin08 studied the elastic properties of austenitic

stainless steel alloy series and measured their lattice parameters. Their

measurements are summarized in Table 6, where E is Young's modulus, G is

shear modulus, B is bulk modulus, v is Poisson's ratio, and V is atomic

volume. Lattice parameter measurements at 4 K are planned in the near

future

.

Table 6. Effect of Alloying on Volume and Elastic Constants.

Alloying

Element

1 &

a dc

1

E

<£

dc

1

G

ac

dc

1

B

dg

dc

1 &

v dc

V(X)/V(Fe)

i i

295 K

I

295 K

1

4 K

1

295 K

1

4 K

I

295 K

H

4 K

i i

295 K 4 K

C 0.218 0.34

-0.78 0.61 -0.76 0.81 -0.93 -0.52 -0.11 -0.90

N 0.240 0.30

Ni -0.003 0.07 — 0.09 •

-0.01 — -0.06 — 0.93

Mn 0.022 -0.30 -0.38 -0.27 -0.37 -0.45 -0.48 -0.12 -0.08 1.04

Mo 0.150 -0.91 -0.14 -1.18 -0.38 0.79 1.30 1.22 1.13 1.32

Cr 0.075 — -

— —— —

— — 1.02

34
At room temperature both C and N show a large aize misfit despite their

small atom size [signified by the V(X)/V(Fe) ratio]; obviously they must

occupy interstitial lattice sites to achieve this Increase. This size

misfit is larger for N than C, although the atom size of N is less than C.

The divergence of E and G trends from B and v trends implies electronic

bonding effects that alter shear modes differently than dilatational modes;

this divergence is found for N and C at 4 K and for Mo at both temperatures.

Furthermore, the change of the sign of the Young's and shear moduli parameters of N and C on cooling from 295 to 4 K implies a low- temperature transition toward increased covalent bonding.

The maximum interaction forces, assuming a - 20, for each solute are

summarized in Table 7. The order of ratios of the interaction forces is

nearly correct when compared to the strengthening contribution listed in

Table 5; only C and N are reversed. However, the relative magnitudes of the

interaction forces of C and N show greater disparity compared to the relative magnitudes of the solute contributions. The interstitial element

interaction forces are smaller than would be expected if one considered

their strengthening contributions. This disparity undoubtedly will be

increased at 4 K when the lattice parameter measurements at 4 K are

obtained: the ratio of strengthening contributions from interstitials

increases relative to substitutional solid- solution elements at 4 K.

This leads to the suggestion that, particularly at low temperatures,

the dislocation-interstitial interaction is distinct from that of the

dislocation-solid- solution interaction. The interaction force ratios could

be brought more in line with the strengthening contributions by increasing

a, the parameter related to the curvature of the interaction potential of

the solvent with the moving dislocations, 83 since the size-misfit parameters

relate directly to the relative strengthening contributions for each type of

atom. These disparities, combined with the linear dependence of strength on

solute interstitial content, lead to the consideration of interstitial

clustering.

Galligan and Goldman03 measured strengthening factors of 2G at room

temperature for presumably split interstitials of Ag in Pb, where Ag has

only slightly lower specific volume than Pb. It is not difficult to

envision that the much smaller C and N interstitials, in similar lattice

positions strengthen by factors of G/10 to G/15 at room temperature and

that this strong, short-range interaction force is strongly temperature

dependent

Table 7. Maximum Interaction Forces* for Solutes in Austenitic

Stainless Steel.

T(K) C N Ni Mn Mo

295 4.43 4.86 0.26 0.52 3.22

*fj " H

2 + a

2 S2

, where a is assumed to equal 20

35
SUMMARY

Austenitic stainless steel structures and strengths have been reviewed.

Emphasis was placed on low- temperature characteristics to include recent

studies that have led to stronger alloys for cryogenic applications. The

stress-strain characteristics of austenitic stainless steels depend on

austenite stability. A region of low work hardening, following initial

plastic deformation, is associated with hep martensite formation and bcc

martensite formation at cross-slip sites. The final work -hardening stage

is associated with bcc martensite formation. All stages are related to

alloying effects on austenite stability.

The interstitials C and N have been shown to contribute more to

strength than that predicted by solution-hardening dislocation theories.

It is argued that C and N must act in clusters to increase solution

hardening. Nitrogen has proven to be the most effective alloying element

for increasing strength at low temperatures.

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41
79 . T. Sakamoto, Y. Nakagava, and I. Yaaauchi." Effect of Mn on the

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in: Advances in Cryogenic Engineering—Materials , vol. 32, Eds.

R. P. Reed and A. F. Clark, Plenum Press, New York (1986), pp. 65-71.

80. Y. Takahashi, K. Yoshida, M. Shimada, E. Tada, R. Miura, and

S. Shimamoto, "Mechanical Evaluation of Nitrogen- Strengthened Stainless

Steels at 4 K," in: Advances in Cryogenic Engineering—Materials , vol. 28, Eds. R. P. Reed and A. F. Clark, Plenum Press, New York

(1986), pp. 73-81.

81. H. nasumoto, K. Suemune, H. Nakajima, and S. Shimamoto, "Development of

High-Strength, High -Manganese Steels for Cryogenic Use," in: Advances

in Cryogenic Engineering—

M

aterials , vol. 30, Eds. A. F. Clark and

R. P. Reed, Plenum Press, New York (1984), pp. 169—176.

82. K. Shibata, Y. Kobiti, Y. Kishimoto, and T. Fujita, "Mechanical

Properties of High Yield Strength High Manganese Steels at Cryogenic

Temperatures," in: Advances in Cryogenic Engineering—Materials . Vol. 30, Eds. R. P. Reed and A. F. Clark, Plenum Press, New York

(1984), pp. 169-176.

83. T. Sakamoto, Y. Nakagawa, I. Yamauchi, T. Zaizen, H. Nakajima, and

S. Shimamoto, "Nitrogen-Containing 25Cr-13Ni Stainless Steel as a

Cryogenic Structural Material," in: Advances in Cryogenic

Engineering -Materials . vol. 30, Eds. A. F. Clark and R. P. Reed,

Plenum Press, New York (1984), pp. 145—152.

84. P. Haasen, "Solution Hardening in f.c.c. Metals," in: Dislocations in

Solids, Chapter 15, Ed. F. R. N. Nabarro, North -Ho Hand, New York

(1976), pp. 155-189.

85. F. R. N. Nabarro, "The Theory of Solution Hardening," Philos. Mag. 21,

613-622 (1977).

86. R. Labusch, "Statistische Theorien der Mischkristallhartung, " Acta

Me tail . 2£, 917-927 (1972).

87. K. Suemune, K. Sugino, H. Matsumoto, H. Nakajima, and S. Shimamoto,

"Improvement of Toughness of a High-Strength, High -Manganese Stainless

Steel for Cryogenic Use," In: Advances in Cryogenic Engineering

Materials . vol. 32, Eds. R. P. Reed and A. F. Clark, Plenum Press,

New York (1986), pp. 51-56.

88. K. Ishikawa, K. Hiraga, T. Ogata, and K. Nagai, "Low Temperature

Properties of High-Manganese-Molybdenum Austenitic Iron Alloys," in:

Austenitic Steels at Low Temperatures . Eds. R. P. Reed and T. Horiuchi,

Plenum Press, New York (1983), pp. 295-309.

89. P. T. Purtscher and R. P. Reed, "Effect of Chemical Composition on the

4 K Mechanical Properties of 316LN-Type Alloys," in: Advances in

Cryogenic Engineering—Materials , vol. 34, Eds. A. F. Clark and R. P.

Reed, Plenum Press, New York (1987), to be published.

42
90 . R. Muira, H. Nakajlma, Y. Takahashi, and K. Yoshida, "32Mn-7Cr

Auatanitlc Steel for Cryogenic Applications," in: Advances in

Cryogenic Engineering Materials , vol. 30, Eds. A. F. Clark and

R. P. Reed, Plenum Press, New York (1984), pp. 245-252.

91. K. Hiraga, K. Ishikava, K. Nagai, and T. Ogata, "Mechanical Properties

of Cold-Rolled and Aged Fe-Ni-Cr-Ti -Austenitic Alloys for Low Temperature Use," in: Advances in Cryogenic Engineering—Materials , vol. 30,

Eds. R. P. Reed and A. F. Clark, Plenum Press, New York (1984),

pp. 203-210.

92. J. M. Galligan and P. D. Goldman, "Metal Interstitial Solid-Solution

S trengthening , " in : Strength of Metals and Alloys. Proceedings of the

Fifth International Conference. Pt. II (1980), pp. 983-988.

93. A. A. Hendrickson and M. E. Fine, "Solid Solution Strengthening of Ag

by Ai," Trans. AIME 221, 967-974 (1961).

94. T. A. Bloom, U. F. Kocks , and P. Nash, "Deformation Behavior of Ni-Mo

Alloys," Acta Metall. 32, 265—277 (1985).

95. Y. Nakada and A. S. Keh, "Solid-Solution Strengthening in Ni-C Alloys,"

Metall. Trans. 2, 441-447 (1971).

96. G. J. denOtter and A. Van den Beukel, "Flow Stress and Activation

Volume of Some Cold-Worked Copper-Based Solid Solutions," Phys. Status

Solidi(a) 22, 785-792 (1975).

97. J. H. Tregilgas and J. M. Galligan, "Hardening from Metal Interstitials

in a Face -Centered Cubic Lattice—Pb-Ag," Scr. Metall. 9, (1975)

pp. 1225-1227.

98. H. M. Ledbetter and M. W. Austin, unpublished measurements, National

Bureau of Standards, Boulder, Colorado (1987).

A3
.
NITROGEN IN AUSTENITIC STAINLESS STEELS

R.P. Reed

National Institute of Standards and Technology

Boulder, Colorado

Nitrogen alloyed in austenitic stainless steels improves

austenite stability, mechanical properties, and corrosion resistance. Research on super-N steels (steels supersaturated with

nitrogen), including processing and properties, is presented.

These new steels rival the latest ferritic steels in strength but

have potentially greater toughness.

INTRODUCTION

For many years nitrogen has been added to austenitic stainless steels

because (1) it stabilizes austenite, thus reducing the tendency for ferrite

formation at high temperatures and martensitic transformation at low temperatures; (2) it increases flow strength more than additions of C or solidsolution alloying elements; (3) it is more soluble than C at intermediate

and high temperatures, thus reducing precipitates; (4) it retards carbide

formation and growth, thus reducing sensitization; (5) it is abundant and

therefore inexpensive. In summary, N imparts to austenitic stainless steels

their serviceable attributes: improved strength and corrosion resistance and

slower high- temperature reactions.

Perhaps the first study of the effects of N additions to steels was by

Adcock1 on Cr and Fe-Cr alloys in 1926. In the 1930s, studies on N additions to Fe—Cr documented their increased strength. In the late 1930s and

early 1940s, the shortage of Ni during the war inspired research on the

possible replacement of Ni with N to stabilize austenite. Replacing Ni with

a combination of N and Mn, another possibility, had the additional advantage

of increasing N solubility. This research in the United States resulted in

new alloys, the AISI-200 series. Gunia and Woodrow2 wrote an excellent

review of the early progress in using N to increase strength, substitute for

Ni, and increase corrosion resistance in austenitic stainless steels.

Zackay used centrifugal casting and pressurized- ingot- solidification

techniques to achieve N contents up to 0.75 wt.% in Fe—Cr—Mn austenitic

steels. 3 * 4 The high- temperature tensile and creep properties of these

alloys were better than those of the current (AISI- 300- series) austenitic

stainless steels. Thus, N- strengthened alloys began to be considered and

developed for high- temperature applications.

45
High-N austenitic steels now in use include the AISI-200 series, N and

HN grades of the AISI -300 series, and specialty steels containing more Cr,

Ni, Mn, or Mo than specified under current AISI designations.

STABILITY

Austenitic steels may solidify as austenite [ face -centered cubic

(fee), 7 ], ferrite [body-centered cubic (bcc) 6 or a ] , or a combination of

these structures. The stability of the austenite structure is complicated.

During cooling or deformation, one or two martensitic products [bcc, a';

hexagonal close-packed (hep), c'] may form. Their formation depends on the

amount of Cr, Ni, Mn, and interstitials (primarily C and N) in the alloy.

When N additions exceed the solubility limits at high temperatures, Cr 2 N

precipitates (with a hep structure) form. Stacking- fault energy depends

primarily on the relative free energies of the 7 and c' structures. The

effects of N addition on these structures are considered.

Solubility

Since high-N contents are usually introduced into austenitic stainless

steels as a gas, it is useful to know the solubility limits of gaseous N

at high temperatures. Solubility depends on pressure; at 1600°C (2900°F),

almost 3 wt.% N has been measured in a molten Fe—20Cr—15Mn alloy at a

pressure of 10 MPa (100 bar). At the same temperature and pressure, about

1 wt.% N can be dissolved in Fe—20Cr—Ni alloys. Results of early work6 at

low nitrogen pressures (< 0.1 MPa or 1 bar) supported Siewert's law: N

solubility is proportional to pressure to the 1/2 power. Siewert's law

breaks down at high pressures: then solubility is proportional to pressure

to a power less than 1/2 . 6 The effects of alloying on N solubility in a

liquid Fe—18Cr—8Ni base at 0.1 MPa (1 bar) and 1600°C (2900 # F) are summarized in Figure l . 6 Most alloying elements increase N solubility, but

the important element, Ni (as well as W, Cu, Sn, Si, and Co), decreases N

solubility. Gaseous -N solubility in austenitic steels is increased by

adding Cr and decreased by adding Ni (see Kikuchi et al . 7 for review).

In the solid, the first N-containing phase to begin precipitating is

usually Cr 2 N. Austenite—Cr 2 N phase -boundary measurements 8 compiled for both

Fe—Cr—Ni and Fe—Cr-Mn austenites and new measurements 7 of the phase boundaries for two austenitic steels, Fe—25Cr—28Ni and Fe—25Cr—20Ni , are shown in

Figure 2. As in the liquid, adding Ni decreases N solubility, and adding Cr

increases N solubility. The austenite—Cr 2 N phase boundary in the Fe—Cr-Mn—

N

system extends to larger values of N, indicating that N is more soluble when

Ni is replaced by Mn. The quaternary Fe—Cr—Ni—N system has been evaluated

thermodynamically . 9

A representative curve for carbide formation in Fe—Cr—Ni—C alloys 10 is

also included in Figure 2. At intermediate temperatures (800—1000°C; 1400—

1800 # F)

,

the solubility of C in solid solution is nearly one -tenth that of

N.

46
- ."sywgpf

Figure 1. Solubility of various alloying elements in Fe—18Cr-8Ni liquid at

0.1 MPa (1 bar) and 1600°C (2900°F). 5

Figure 2. Phase boundaries between austenite and Cr 2 N or carbides (for

Fe—Cr-Ni—C alloys) with addition of N or C. 7,9 » 10 ].

47
Solidification

The relative amounts of austenite and ferrite that form during solidification are critically dependent on alloying. The liquidus and solidus surfaces of the Fe—Cr—Ni ternary have been determined: The eutectic trough

extends from near the Fe comer (76Fe—lONi—14Cr) to a ternary eutectic (about

8Fe-43Ni-49Cr )

.

1

1

Primary austenite solidifies when the composition is on the

Ni side of the eutectic liquidus. Nearer the eutectic liquidus on the Ni

side, limited ferrite forms, mainly in the dendritic cell boundaries. On the

Cr-rich side of the liquidus line, primary ferrite solidifies. Upon cooling,

however, ferrite transforms to austenite in compositions typical of austenitic

stainless steels. Although few quantitative experimental studies of N effects

on primary solidification modes have been reported, the stabilization of austenitic structure by N addition is recognized.

For estimates of ferrite and austenite in the weldments of austenitic

stainless steels, solidification equivalent coefficients for Cr and Ni have

been empirically developed (see Reed 10 for review). Expressions relating

these coefficients to chemical composition are plotted on a diagram similar to

the one originally developed by Schaeffler . 12 The N contribution is included

as part of the Ni (austenite stability) term. Currently, on a weight-percent

basis, N is considered to be 20 times more effective than Ni in stabilizing

austenite; that is, the solidification equivalent coefficient for N is 20

times the Ni coefficient . 13

Ferrite

Ferrite may form during either solidification or high- temperature treatment, depending on alloy content. Occasionally small amounts (usually <10%)

are present in rapidly cooled, lean, austenitic (AISI-300-series) stainless

steel alloys (those whose primary solidification mode is ferrite), such as

weldments. This bcc phase is conventionally called 6 -ferrite when it forms at

temperatures above the 7 - loop or range, and a- ferrite when it forms at lower

temperatures. Most austenitic stainless steels have sufficient stabilizing

elements (such as Ni, Mn, C, and N) to lower the temperature of the 7 - loop

enough to prevent the diffusion-controlled 7

-* a transformation. In wrought

alloys, the a-ferrite retained from solidification is typically transformed to

austenite during forging or hot rolling.

Figure 3 shows the effect of N on the boundary of the austenite and austenite + ferrite phases for an Fe-18Cr-base alloy at various temperatures . 14

Below 1200°C (2200 # F) , adding N only mildly suppresses ferrite formation

(1 wt.% N * 20 wt.% Ni when the slope of the phase boundary is considered).

Above 1200°C, N strongly stabilizes the austenitic structure in alloys containing 7.5 wt.% Ni (1 wt.% N “ 200 wt.% Ni) . Thus, the lower temperature

N:Ni equivalency ratio (1:20) corresponds to the weld-metal ratio discussed

above, indicating sufficient time and solute mobility below 1200°C to establish near-equilibrium conditions during weld-metal solidification.

48
Figure 3. Phase boundaries separating austenite, austenite + ferrite, and

austenite + Cr 2 N regions as a function of N and Ni additions to an

Fe—18Cr base alloy, plotted for various temperatures. 14

Martensitic Transformation

The addition of N increases austenite stability with respect to martensitic transformation in both bcc and hep structures. The suppression of a'

martensite by N on cooling has long been recognized, and expressions to

calculate the martensite start temperature, Tms , were developed (see, for

example, the review by Reed 10 ). In these expressions, the contributions of

N and C have usually been found to be the same: 1 wt.% N or C lowers Tms by

roughly 1500 # C (2700°F). In most Tms expressions, N contributes about 25

times more than Ni, close to the ratios of 20 for equilibrium and weldmentferrite formation below 1200°C (2200°F), which were discussed earlier. In expressions for calculation of the highest temperature for strain- induced

a' martensite, Tm(j.

10 the contributions of N and C are also equivalent: 1 wt.%

of interstitial increases Tm<j about 250 °C (450 °F) , approximately

10 times that of the common metal alloying elements (Cr, Ni, Mn)

The effect of N addition on c

' -martensite formation in austenitic steels

has not been previously characterized. Our recent room- temperature measurements 16 of Fe—18Cr—lONi alloy specimens plastically deformed to strains of

0.05 and 0.10 at 76 K (— 323°F), shown in Figure 4, were obtained by using

integrated x-ray diffraction peaks of «' martensite. Adding both

C (0.038 — 0.112 wt.%) and N (0.042 — 0.264 wt.%) increases austenite

stability with respect to strain- induced e' martensite.

49
>

w

c

03 JZ>

CO

v_ ><

TD

0)

CO

i_O)

0)c

<13 >

OJ

d) QC

Figure 4. Relative integrated peak intensity, averaged from the (100)

and (101) hep structure peaks, following plastic deformation to

strains of 0.05 and 0.10 at 76 K (- 321 °F); x-ray analysis at

295 K (72 # F)

.

1

6

Stackjng-£aul t gnergy

Stacking- fault energy is related to the free-energy difference between

the fee principal structure and the hep structure. Lower stacking- fault

energy leads to wider partial-dislocation separation, more planar dislocation

arrays, and reduced cross slip.

The effects of additions of C and N (C + N < 0.364 wt.%, N ^ 0.264 wt.%)

on the stacking- fault energy of an Fe—18Cr—lONi alloy were examined with x- rays. 18 Within data scatter, the stacking- fault energy can be considered to

be independent of C + N content. Using transmission electron microscopy,

values of the stacking -fault energy of 24 ± 3 mJ/m2 and 23 ± 5 mJ/m2 were

measured for a Fe—17Cr—12Ni—2Mo-0 . 02C alloy with N contents of 0.08 and 0.25

wt.%, respectively. 17 Weak-beam microscopy measurements showed that N in

excess of 0.24 wt.% (to 0.52 wt.%) had little effect on the stacking- fault

energy of an Fe—19Cr-7Ni—8Mn—0 . 03C alloy, but that at 0.21 wt.% N, the

stacking -fault energy was considerably higher. 18 All three sets of measurements may be consistent with either of the following possibilities: (1) The

stacking- fault energy is independent of N contents in Fe—Cr—Ni alloys; in

Fe—Cr—Ni-Mn alloys, there is a transition. (2) Both alloy systems undergo a

transition in the stacking- fault energy with N content, but in the Fe—Cr-Ni

alloys, the transition occurs at higher N content, and the N-solubility limits

may preclude measurement. However, it is discomforting to realize that these

50
"transition" premises depend exclusively on one measurement (of the Fe—

19Cr—7Ni-8Mnr-0.03C—0. 21N alloy 18

); all other results indicate no N-composition

dependence of the stacking- fault energy. The constant (or reduced) stackingfault energy and the suppression of the 7 -* e

transformation with increased N

content are conflicting trends since both are expected to reflect the

dependence of the free -energy difference between the fee and hep structures on

N content.

CORROSION RESISTANCE

The increasing demands of the petrochemical industry for higher strength,

corrosion-resistant alloys have led to consideration of N-alloyed austenitic

stainless steels. Additions to high-Cr (> 20 wt.%) austenitic steels containing Mo and about 0.2 wt.% N significantly increase the resistance to corrosion

from acids (Figure 5 ). 19,20 Nitrogen actively assists in the formation of a

preliminary film, and apparently, subsequent passive films contain more Cr and

Mo if N is present . 21 When N is added to common austenitic steels (such as

AISI 304, 316, and 347), it has less effect on corrosion resistance to

acids—and sometimes even a slightly adverse effect . 22

Pitting corrosion, which usually occurs in the presence of halide (Cl,

Br) ions, represents localized breakdown of the passive film. Figure 5 shows

that adding N decreases the pitting corrosion rate from FeCl s -HCl attack.

With these reduced rates, an increase in the effective potential is needed to

initiate pitting corrosion. The resistance to pitting corresponds to l[Cr] +

3.3[Mo] + x[N]

,

where [Cr] , [Mo], and [N] represent the total wt.% of each

element and values of x range from 13 to 30. 22,23 Plots of this expression

versus effective pitting potential 22 reach a plateau maximum at about a value

of 33 when x is 16.

Truman22 summarized various studies relating N content to stress corrosion and corrosion fatigue. Most authors concluded that increased N levels

reduce the effects of stress -enhanced corrosion, but a few disagreed.

Nitrogen content usually reduces the deleterious effects of sensitization

(precipitation or micros truetural solute segregation) on corrosion resistance

However, when the N content exceeds its solubility limits, Cr 2 N will precipitate and deplete the matrix of Cr; this results in reduced passivity.

Carbides precipitate at ferrite—austenite interfaces, at grain boundaries, at noncoherent and coherent twin boundaries, and finally, within grains

(in order of probability). As the precipitation site changes from high- to

low-energy boundaries, the temperature decreases and the time for nucleation

increases. Increased C content increases precipitation temperature and

decreases time; increased N content has the reverse effect, suppressing

carbide formation . 19 * 24 The M23 C 8 carbide has a complex fee structure with 92

metallic atoms and 24 smaller (C, N, or B) atoms. The common plane of the

carbide—austenite interface has been identified as {111 }. 26 Alloy additions

to the M23 C 6 structure, such as N, tend to decrease its lattice parameter and

to increase the amount of interface mismatch. This should retard the growth

kinetics

51
STRENGTH

Because N provides excellent strength with very little increase in production cost, it has been used in the development of strong, tough, lowtemperature alloys. Table I summarizes the strength increases calculated

(in most cases using regression analyses) from controlled experimental data

sets of many studies; 26-36 various techniques or assumptions were used to determine the dependence of the power of the N content ([N]) for calculating the

yield strength. Some analyses assumed a linear fit, 28-31 ’ 34 ’ 35 some

used [ N ] ^ , 2 6 * 32 ,33 an(j some compared various powers of the N content. 26 ’ 27

From comparisons for the same data, powers less than 1 usually result in

slightly better data fits, but these trends are usually within 2 standard

deviations of each other. Therefore, there is no conclusive evidence of the

dependence of the yield strength on the power of [N]

.

Table I. Nitrogen Contribution to Yield Strength from Various Reports

Authors

(Date)

Base Alloy N Contribution to

MPa/wt .

%

4 K (—452 # F)

Yield Strength,

(ksi/wt . %)

295 K (72°F)

Reed, Simon (1989) 26 Fe-19Cr-10Ni 2400 (348) 500 (73)

Reed, Purtscher,

Delgado (1988) 27

Fe—Cr—Mn 2140 (311) 680 (99)

Werner, Uggowitzer,

Speidel (1987) 26

Fe-Cr-Mn 580 (84)*

Simon, Reed (1986) 27 Fe-Cr-Ni

Fe—Cr—Ni-Mo

3190 (463)

2780 (403)

Yamamoto , Yamagarni

,

2 8

Ouchi (1986)

Fe—Cr—Mn 2940 (427) 570 (83)

Reed, Simon (1984) 27 Fe-Cr-Ni 3400 (493) 420 (61)

Sakamoto et al. (1984) 30 Fe-Cr-Ni 2500 (363) 690 (100)

Takahashi et al. (1982) 31 Fe-Cr-Ni-Mo 2910 (422)

Fujikura, Kato (1978) 32 Fe-Cr-Mn 510 (74)

Norstrcm (1977) 33 Fe—18Cr—14Ni—3Mo 310 (45)*

Irvine, Gladman,

Pickering (1969) 34

Fe-Cr-Ni 440 (64)

I rvine , Llewe 1 lyn

,

Pickering (1961) 35

Fe—18Cr—lONi 710 (103)

An interactive contribution from N grain size was also found.

52
Figure 5. Dependence of corrosion rate on N content for both an acidic

(H 2 S04 , HC1) and crevice (FeCl s ) environment at different

temperatures

1

0

The strengthening contribution from interstitial N addition is very

dependent on temperature. Figure 6 presents average measurements at 295,

77, and 4 K (72, -321, and -452 # F) for Fe-Cr-Ni alloys with N contents

within the solubility limits. At 295 K, the increase in yield strength

is approximately 50 MPa (7.3 ksi)/0.1 wt.% N addition. At 4 K, this

strengthening has increased to about 250 MPa (36 ksi)/0.1 wt.% N—a factor

of 5.

53
Figure 6. Dependence of the tensile yield strength (0.2% offset) on N

content of Fe—Cr—Ni-base austenitic alloys for selected low

temperatures

.

2 ®» 30 > 31 > 3 ®» 37

Figure 7 shows the effects of common alloying elements on the yield

strength of Fe—Cr—Ni-base austenitic stainless steels at 295 K. 3S The largest increases in yield strength are supplied by the smallest elements (N, C,

and B) that occupy interstitial sites. Nitrogen strengthens steel more than

C: At 295 K, N increases strength about 20% more than C, but at 4 K, N

increases strength 300% more than C. 26 Ferrite -forming elements, elements

whose atomic sizes are much different from those of the Fe—Cr—Ni solvent,

are next in order of strengthening contributions. Elements whose atomic

sizes closely match the average size of the Fe—Cr-Ni matrix contribute the

least

.

54
CD

C

03

O

03

Q_

240

200

160

JO

O) 120

c

<D 80

cn 40

2

<D 0

>

-40

(

substitutional,

ferrite forming

—Co, Mn ) \ ... .

1 ( substitutional,

(austenite forming

4.0 6.0 8.0 10.0 12.0

Addition, at. %

16.0

Figure 7. Contributions of selected alloying elements to 295 -K tensile

yield strength for Fe—Cr—Ni-base austenitic stainless steels. 35

The large increases in flow strength that result from N alloying remain a puzzle. Proposals that account for the discrepancies from solutionhardening theories include long-range ordering; local ordering between N and

Cr, Mn, or Mo; split interstitial sites; "microprecipitates" of nitrites;

clustering; and Mn—N effective "molecules." Therefore, identification of

the actual occupancy sites and randomness of distribution for N atoms will

significantly improve our understanding of interstitial strengthening in fee

alloys. One piece of evidence is the frequent observation that dislocations

tend toward a linear alignment at high-N contents and toward a cellular

structure at low-N contents. The relative insensitivity to stacking- fault

energy at compositions below about 0.24 wt.% N (see discussion of stackingfault energy) implies that ordering may explain the change of dislocation

geometry.

CREEP AND FATIGUE

For high- temperature applications, such as in steam and nuclear power

plants, N alloying of austenitic stainless steels usually prolongs operational lifetime. As Figure 8 illustrates, adding N also enables the attainment of higher stress for a desired operational lifetime. 38 ” 40 Creep

ductility is not sacrificed, 41 and at times, is improved. 42 * 43 In contrast,

aduing C lowers creep ductility. 41 Figure 8 shows that very small changes

in interstitial content substantially improve creep resistance.

55
Figure 8. Effect of N and C content on creep strength to achieve a 100,000-

h lifetime for Fe—19Cr—lONi base alloys (with N contents from

0.03 to 0.14, C contents from 0.02 to 0.10) at holding temperatures of 650°C (1200°F) and 730°C (1350°F). Band represents more

than 40 data points (650°C data from Goodell, Cullen, and

Freeman; 38 730°C data from Hoke review39 ).

There is evidence that the improved creep properties resulting from N

addition are associated with N solid-solution strengthening. 44 When a series of alloys with very low C was used to preclude the influence of carbide

formation, the minimum creep rates decreased with increasing N content from

800 to 1000°C (1500 to 1850°F) (Figure 9). Also, the decrease of minimum

creep rate at 800°C was independent of C content, as illustrated by the

dashed curve in Figure 9. In this study, intergranular Cr 2 N precipitates at

lower temperatures and at long times to rupture were observed. The minimum

creep rate decreased with increasing N content in all cases where the N

remained in solid solution.

At temperatures higher than about 800°C (1500°F), the distinctions

between C and N become less. In fact, at 900°C (1650 # F), the creep properties of AISI 316 were reported to be superior to those of AISI 316LN. 42 At

lower temperatures (< 800°C) , the reduced amount of N in solid solution

seems to be dependent on the precipitation of Cr 2 N. At higher temperatures

(900; 1000°C) , the more similar creep behavior of low- and high-N austenitic

steels may be attributed to the participation of N in carbide formation.

The coarse carbonitrides would, of course, remove N from solid solution and

reduce its strengthening effect.

56
Figure 9. Ratio of minimum creep rate, € m^n to a N-free alloy and a 0.005-

wt.%-N alloy with 0.008 wt.% C as a function of N content at a

stress level of 48 MPa (7 ksi) ; solid lines. The dashed line

represents data at 800°C (1500°F) and a stress level of 98 MPa

(14 ksi) for all alloys with varying C + N content. 44

The fatigue -crack- growth rates are affected by strain- induced martensite. Martensite formation in the plastic zone about the crack tip shields

the crack, lowering the stress intensity at the tip of the crack. This is

illustrated by the low- temperature values 45 for a series of AISI 304 alloys

with varying N and C contents (Figure 10) . At 295 K, high-strength alloys

(those with high C + N contents) had lower fatigue -crack- growth rates. At

low temperatures, the crack-growth rates of the high-strength alloys were

the same or higher than the 295-K rates, but the crack-growth rates of the

low- strength steels (low C + N content) were considerably lower. This

reduction correlates very well with the increase of strain- induced a'

martensite in the process zone at the crack tip.

The effects of N on fatigue behavior were recently reviewed. 46

Nitrogen improves the fatigue life of austenitic steels; most studies

attributed this improvement to increased resistance to strain localization

from increased planar slip. Resistance to strain localization reduces the

formation of persistent slip bands and increases slip reversibility.

57
13

75

a

af

75cc

$

o

<5

o

03 O

<D

ZJ _gj

75

LL

Fe-18Cr-10Ni Stainless Steels

Figure 10. Fatigue -crack- growth rates at low temperatures for Fe—

18Cr—lONi base alloy with varying C + N contents that are

specified in wt.% with each curve. 45

SUPER-NITROGEN STEELS

Processing

The development of pressurized electroslag remelting as an industrial

process has enabled the production of super-N austenitic stainless steels

(steels supersaturated with N) . These new steels contain up to 1 wt.% N, in

excess of the N solubility at atmospheric pressure. Super-N steels are now

being produced for generator retaining rings, and they are expected to be

used in power generation, ship building, chemical and cryogenic processes,

and pressure vessels. Processes that were proposed previously for producing

super-N steels have not been amenable to large-scale industrial production

(pressurized induction furnaces, 6 centrifugal casting4

), have tended to

produce inhomogeneous ingots (plasma-arc remelting47

), or have not been able

to achieve high enough N levels (ingot gasification4 ).

58
The use of pressurized electroslag remelting is extensively discussed

by Stein et al. in their review of some of the super-N- steel processes. 48

In both plasma- arc remelting and pressurized induction melting, N is

introduced as a gas, but in the pressurized electroslag remelting process,

the N is introduced to the liquid-phase alloy in granulated form within the

slag system. Granulated additives with high N contents must be used;

currently, silicon nitrides are favored. 48

To obtain the desired chemical composition and homogeneity, it is

necessary to consider electrode and additive chemistry, speed of motion

through the slag and liquid, bath convection, temperature, and N pressure.

Double remelting is sometimes necessary to obtain adequate chemical homogeneity. Currently, the most advanced pressurized-electroslag-remelting

facility is capable of producing 20 -ton ingots up to 1000 mm (39 in) in

diameter. 48 It uses a stationary copper mold that is cooled with pressurized water, has a movable hood, and is capable of a maximum pressure of

4.2 MPa (42 bar). In contrast to normal electroslag-remelting practice, the

pressurized system requires more automation because the progress of the

remelting is not visible. Thus, a reliable sensor system and knowledge of

the effects of pressure, temperature, and speed variations are necessary for

efficient industrial application.

Super-N steels are more difficult to hot-work. The range of hotworking temperatures and amount of hot working are restricted, limiting the

size and shape of the forgings that can be achieved within given mill

facilities

Holzgruber48 maintains that super-N steels are best produced on a large

scale by using N-gas purging of liquid metal. In this process, electroslag

heating with either a consumable or nonconsumable electrode produces a

liquid-metal alloy with a constant chemical composition that is unaffected

when N is added. Gaseous N is added to this base metal from the bottom; the

liquid metal and slag pools are held under pressure to control the N

content. An important advantage of this system is that any normal casting

technique can be used, since the entire process can be performed in a ladle

like container. For large-scale production of ingots, a pressurized casting

chamber, featuring a porous bottom plug in the liquid-metal container and a

launder with pressure -gate valves through which the molten metal pours into

the ingot mold, is suggested.

SUPER-NITROGEN STEELS

Properties

A completely new set of property data is being developed from super-N

steels produced by the pressurized-electroslag-remelting process. In

Figure 6, the Fe-Cr—Ni-base data have been extended to N contents from

0.4 to 1.0 wt.%. 36 ’ 37 The linear trend of tensile yield strength versus N

content at room temperature to approximately 1 wt.% encourages the use of

alloys with N contents in excess of their solubility limits. Yield

strengths to almost 1000 MPa (145 ksi) at room temperature are attainable in

alloys with no cold or hot working and with a single phase.

59
The most significant information, though, is that super-N steels are

comparable to, if not better than, contemporary ferritic steels in terms of

strength and toughness (Figure ll). 36 * 50 * 51 The higher strength was

attained by cold-working at room temperature. The success of acquiring very

high strengths in austenitic steels is particularly encouraging because, in

trast to the maraging and dual-phase steels and 9Ni-Co alloys, the development of super-N steels did not have the benefit of extensive research

programs to optimize strength—toughness behavior through chemical composition and thermomechanical processing.

Other property data on super-N austenitic steels have resulted from

studies of welding, 48 effects of grain size on tensile yield strength, 28

the temperature dependence of the tensile flow strength, 8 * 87 strain hardening, 8 * 37 effects of annealing time and temperature on fracture toughness

and Cr2 N precipitation, 8 corrosion resistance, 38 * 52 effects of cold work

on tensile yield strength, 36 * 50 * 52 and reversed- cycle fatigue at room

temperature

.

53

Yield Strength, ksi

100 200 300 400

500

1l/l -X

400 ~

o

)Z

T3

300 <°o

)Z

cn" <o 200 c

r.O)

3

|2

0) 100 5

o

U.

~0 500 1000 1500 2000 2500 3000

Yield Strength, MPa

Figure 11. Room- temperature fracture toughness and yield strength data to

compare super-N steels with ferritic steels. Fracture toughness

data either from linear elastic tests or from J- Integral tests,

converted to linear elastic. 36 * 50 * 51

60
CONCLUSIONS

Additions of N to austenitic stainless steels increase austenite stability, strength, and corrosion resistance and improve performance. These

improvements enable more extensive use of high-N steels at both low and high

temperatures. The very recently developed super-N austenitic steels have

potential for applications where very high performance steels are required.

REFERENCES

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Alloys," J. Iron Steel Inst. 114 (1926) 119-126.

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of Chromium-Nickel Stainless Steels," J. Mater. 5 (1970) 413—430.

3. V.F. Zackay, J.F. Carlson, and P.L. , Jackson, "High Nitrogen Austenitic

Steels," Trans. Amer. Soc. Met. 48 (1956) 509-525.

4. V.F. Zackay, E.R. Morgan, and J.C. Shyne , "Production of High Nitrogen

Steels," J . Met. 8 (1956) 216.

5. W.M. Small and R.D. Pehlke, "The Effect of Alloying Elements on the

Solubility of Nitrogen in Liquid Iron-Chromium-Nickel Alloys," Trans.

AIME 242 (1968) 2501-2505.

6. H. Feichtinger, A. Satir-Kolorz , and Z. Xiao-Hong, "Solubility of

Nitrogen in Solid and Liquid Iron Alloys with Special Regard to the

Melting Range," accepted for publication in Proc. "HNS -88" (Lille,

France) ed. J. Foct and A. Hendry (London: Institute of Metals, 1989).

7. M. Kikuchi, M. Kajihara, and K. Frisk, "Solubility of Nitrogen in Austenitic Stainless Steels," accepted for publication in Proc. "HNS -88"

(Lille, France) ed. J. Foct and A. Hendry (London: Institute of Metals,

1989).

8. M.A. Harzenmoser and P.J. Uggowitzer, "Neue aufgesteckte austenitischrostfreie StAhle und Duplexs table , " Modem St&hle, vol. 1, ed. P.J.

Uggowitzer and M.O. Speidel (Zurich: Schweizerischen Akademie der

Werkstoffwissenschaften, ETH, 1987), 219-247.

9. K. Frisk and M. Hillert, "Thermodynamics of the Fe—Cr—Ni—N System,"

accepted for publication in Proc. "HNS -88" (Lille, France) ed. J. Foct

and A. Hendry (London: Institute of Metals, 1989).

10. R.P. Reed, "Austenitic Stainless Steels with Emphasis on Strength at

Low Temperatures," Alloying (Metals Park, OH: ASM International, 1988),

225-256.

11. G.R. Speich, Source Book on Stainless Steels (Metals Park, OH: ASM

International, 1976), 424-426.

12. A.L. Schaeffler, "Constitution Diagram for Stainless Steel Weld

Metals," Met. Prog. 56 (1949) 680.

13. T.A. Siewert, C.N. McCowan, and D.L. Olson, "Ferrite Number Prediction

to 100 FN in Stainless Steel Weld Metal," Weld. J. 67 (1988) 289s—298s.

14. T. Masumoto and Y. Imai, "Structural Diagrams and Tensile Properties of

the 18%Cr-Fe—Ni—N Quaternary System Alloys," J. Jap. Inst. Met. 33

(1969) 1364-1371.

15. R.P. Reed, unpublished data, National Institute of Standards and Technology, Boulder, Colorado (1988).

61
16 . H.M. Ledbetter and M.W. Austin, "Stacking -Fault Energies in 304 -Type

Stainless Steels: Effects of Interstitial Carbon and Nitrogen," in

Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures — VIII , ed. R.P. Reed (NBSIR 85-3025, National Institute

of Standards and Technology, 1985), 271-294.

17. R. Taillard and J. Foct, "Mechanisms of the Action of Nitrogen Interstitials upon the Low Cycle Fatigue Behavior of 316 Stainless Steel,"

accepted for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct

and A. Hendry (London: Institute of Metals, 1989).

18. R.E. Stoltz and J.B. VanderSande, "The Effect of Nitrogen on Stacking

Fault Energy of Fe—Ni-Cr-Mn Steels," Metall. Trans. 11, (A) (1980)

1033-1037.

19. T. Sakamoto et al., "High Corrosion Resistent Nitrogen-Containing

Stainless Steels for Use by the Chemical Industry," Alloys for the

Eighties, ed. R.Q. Barr (Greenwich, CT: AMAX, 1980), 269-279.

20. J.R. Kerns, "The Effect of Nitrogen on the Corrosion Resistance of Austenitic Stainless Alloys Containing Molybdenum," New Developments in

Stainless Steel Technology , ed. R.A. Lula (Metals Park, OH: ASM International, 1985), 117-127.

21. C.R. Clayton and K.G. Martin, "Evidence of Anodic Segregation of Nitrogen in High Nitrogen Stainless Steels and Its Influence on Passivity,"

accepted for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct

and A. Hendry (London: Institute of Metals, 1989).

22. J.E. Truman, "Effects of Nitrogen Alloying on Corrosion Behavior of

High Alloy Steels," accepted for publication in Proc. "HNS-88" (Lille,

France) ed. J. Foct and A. Hendry (London: Institute of Metals, 1989).

23. P. GQmpel and T. Ladwein, "The Effect of Nitrogen on MechanicalTechnological and Corrosion Properties of Stainless Steels," accepted

for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct and A.

Hendry (London: Institute of Metals, 1989).

24. H. Their, A. Baumel, and E. Schmidtmann, "Einfluss von Stickstoff auf

des Ausscheidungsverhalten des Stahles x5CrNiMol713 , " Arch. Eisenhhttenw. 40 (1969) 333-339.

25. M.H. Lewis and B. Hattersley, "Precipitation of M2 sC 6 in Austenitic

Steels," Acta Metall. 13 (1965) 1159-1168.

26. R.P. Reed and N.J. Simon, "Nitrogen Strengthening of Austenitic Stainless Steels at Low Temperatures," accepted for publication in Proc.

"HNS-88" (Lille, France) ed. J. Foct and A. Hendry (London: Institute

of Metals , 1989) . 27. R.P. Reed, P.T. Purtscher, and L.A. Delgado, "Low-Temperature Properties of High -Manganese Austenitic Steels," High -Manganese Austenitic

Steels, ed. R.A. Lula (Metals Park, OH: ASM International, 1988),

13-22.

28. E. Werner, P. J. Uggowitzer, and M. 0. Speidel, "Mechanical Properties

and Aging Behavior of Nitrogen Alloyed Austenitic Steels," Mechanical

Behaviour of Materials — V, (New York: Pergamon, 1987).

29. S. Yamamoto, N. Yamagami, and C. Ouchi, "Effect of Metallurgical Variables on Strength and Toughness of Mn—Cr and Ni—Cr Stainless Steels at

4.2 K," Advances in Cryogenic Engineering - Materials , vol. 32, ed.

R.P. Reed and A.F. Clark (New York, NY: Plenum, 1986), 57—64.

30. T. Sakamoto et al., "Nitrogen- Containing 25Cr-13Ni Stainless Steel as

62
a Cryogenic Structural Material," Advances in Cryogenic Engineering -

Materials

,

vol. 30, ed. A.F. Clark and R.P. Reed (New York, NY: Plenum,

1984), 137—144.

31.

Y. Takahaahi et al., "Mechanical Evaluation of Nitrogen- Strengthened

Stainleaa Steels at 4 K," Advances in Cryogenic Engineering — Materials, vol. 28, ed. R.P. Reed and A.F. Clark (New York, NY: Plenum, 1982),

73-81.

32. M. Fujikura and T. Kato, "Effects of Carbon, Nitrogen and Nickel Content

on the Low Temperature Impact Value of High Manganese Austenitic Steel,"

J. Iron Steel Inst. Jap. 64 (1978) 97-106.

33. L.A. Norstrdm, "The Influence of Nitrogen and Grain Size on Yield

Strength in Type AISI 316L Austenitic Stainless Steel," Met. Sci. 11

(1977) 208-212.

34. K.J. Irvine, T. Gladman, and F.B. Pickering, "The Strength of Austenitic

Steels," J. Iron Steel Inst. 207 (1969) 1017—1028.

35. K.J. Irvine, D.T. Llewellyn, and F.B. Pickering, "High-Strength Austenitic Stainless Steels," J. Iron Steel Inst. 199 (1961) 153-175.

36. G. Stein, J. Menzel, and H. D6rr, "Mdglichkeiten zur Herstellung von

Schmiedestrucken mit hohen Stickstoffgehalten in der Desu-Anlage ,

Modeme St&hle, Ergebnisse der Werkstoff—Forschung , vol. 1 (Zurich:

Schweizerische Akademie der Werstoffwissenschaften, 1987), 181—193.

37. P.J. Uggowitzer and M. Harzenmoser, "Strengthening of Austenitic Stainless Steels by Nitrogen, accepted for publication in Proc. "HNS-88"

(Lille, France) ed. J. Foct and A. Hendry (London: Institute of Metals,

1989).

38. P.D. Goode 11, T.M. Cullen, and J.V. Freeman, "The Influence of Nitrogen

and Certain Other Elements on the Creep-Rupture Properties of Wholly

Austenitic Type 304 Steel," J. Basic Eng. (Trans. Amer. Soc. Met. E) 89

(1967) 517-524.

39. J.H. Hoke, "Mechanical Properties of Stainless Steel at Elevated Temperatures," Handbook of Stainless Steels, ed. D. Peckner and I.M. Bernstein

(New York, NY: McGraw-Hill, 1977), 21.1-21.20.

40. Y. Kawabe, R. Nakagawa, and T. Makoyama, "Effect of Nitrogen and Molybdenum on the High Temperature Strength of 18Cr—12Ni-0.02C Austenitic

Steel," Trans Iron Steel Inst. Jap. 8 (1968) 353-362.

41. T. Nakazawa et al., "Effects of Nitrogen and Carbon on Creep Properties

of Type 316 Stainless Steels," accepted for publication in Proc. "HNS-88"

(Lille, France) ed. J. Foct and A. Hendry (London: Institute of Metals,

1989).

42. J.K. Solberg, "The Influence of Carbon and Nitrogen on the High Temperature Creep Properties of AISI 316 Austenitic Stainless Steel," Mater.

Sci. Eng. 55 (1982) 39-44.

43. J. Sobotka et al., "Creep Ductility and Creep Fracture of Non-Stabilized

CrNiMo Steels with Nitrogen Additions," accepted for publication in Proc.

"HNS-88" (Lille, France) ed. J. Foct and A. Hendry (London: Institute of

Metals, 1989).

44. T. Matsuo et al., "Effect of Nitrogen on Creep Deformation of 25Cr—28Ni

Austenitic steels - Solid Solution Strengthening Due to Nitrogen, accepted for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct and

A. Hendry (London: Institute of Metals, 1989).

45. R.L. Tobler and R.P. Reed, "Interstitial Carbon and Nitrogen Effects on

63
the Cryogenic Fatigue Crack Growth of AISI 304 Stainless Steel," J. Test.

Eval . 12 (1984) 364-370.

46. S. Degallaix, J. I. Dickson, and J. Foot, "The Effect of Nitrogen on the

Fatigue and Creep-Fatigue Behavior of Austenitic Stainless Steels," accepted for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct and

A. Hendry (London: Institute of Metals, 1989).

47. G. Torkhov et al., "Development of Melting and Thermomechanical Processing Parameters for a High-Nitrogen Stainless Steel Prepared by Plasma-Arc

Remelting," J. Met. 30 (1978) 20-27.

48. G. Stein, J. Menzel, and H. D6rr, "Industrial Manufacture of Massively

Nitrogen-Alloyed Steels," accepted for publication in Proc. "HNS-88"

(Lille, France) ed. J. Foct and A. Hendry (London: Institute of Metals,

1989).

49. W. Holzgruber, "Process Technology for High Nitrogen Steels," accepted

for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct and

A. Hendry (London: Institute of Metals, 1989).

50. P.J. Uggowitzer, M.O. Speidel, and E. Werner, "Fracture Toughness of Cold

Worked Austenitic Steels," Mechanical Behaviour of Materials — V (New

York: Pergamon, 1987), 401-410.

51. E.B. Kula, Army Materials Technology Laboratory, Watertown, Massachusetts; private communication with author, 1988.

52. M.O. Speidel, "Properties and Applications of High-Nitrogen Steels,"

accepted for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct

and A. Hendry (London: Institute of Metals, 1989).

53. J. Menzel, G. Stein, and P. Dahlmann, "Massively Nitrogen-Alloyed Bolt

Materials for High-Strength and High -Temperature Applications," accepted

for publication in Proc. "HNS-88" (Lille, France) ed. J. Foct and

A. Hendry (London: Institute of Metals, 1989).

64
EFFECT OF NITROGEN AND CARBON ON FCC-HCP STABILITY IN AUSTENITIC STEELS*

R.P. Reed and M.W. Austin

National Institute of Standards and Technology

Boulder, Colorado

INTRODUCTION

High-nitrogen austenitic steels have recently been developed for lowtemperature use.

1 For a long time it has been recognized that both nitrogen

and carbon increase austenite stability with respect to formation of delta

ferrite at high temperatures and body-centered cubic (bcc) martensite at low

temperatures. 1 ’ 2 Yet, the effects of nitrogen and carbon on austenite (fee)

stability relative to hexagonal close-packed (hep) martensite and the

stacking-fault energy have not been adequately clarified.

MATERIALS AND PROCEDURES

The austenitic alloys used for this program are listed in Table 1. The

solid-solution solutes Cr, Ni, and Mn were held constant; the interstitial

solutes C and N were varied. Grain sizes ranged from 80 to 100 um. These

alloys are a subset of a series of laboratory heats of nine alloys that were

previously used to measure strength, toughness, lattice parameters, and

elastic moduli from 4 to 295 K. 1

TABLE 1. Chemical Composition (wt. I)

Alloy C N Cr Ni Mn

1 0.038 0.042 18.3 10.1 1.57

3 0.112 0.044 18.3 10.1 1.51

5 0.076 0.135 18.0 9.9 1.50

9 0.095 0.264 20.3 10.0 1.42

The alloys also contained 0.02 P, 0.02 S, 0.5 Si,

0.2 Mo, and 0.2 Cu.

The tensile specimens were flat with a 38 -mm gage length that was

6.4 mm wide and 3.2 mm thick. — These specimens were incrementally deformed

in tension at 76 K to plastic strains of 0.05 and 0.10. Following the first

Submitted to Scripta Metallurgica

65
strain increment, the specimens were warmed to roc* temperature, and x-ray

diffraction peaks were measured. Then the specimen ?ere returned to 76 K

for the second strain increment.

A horizontal Bragg-Brentano type of x-ray diffractometer was used with

Cu radiation, a 0.3-mm receiving slit, and a Ge solid-state detector set for

K-alpha radiation. Specimens were scanned for 12 s per 0.02° step at room

temperature. Data files were constructed by subtracting the background,

correcting for the Lorentz polarization, and deleting the Cu K-alpha-2 contribution with a Rachinger algorithm.

To analyze the fee structure, diffraction peaks (111), and (200). were

used; to analyze the hep strain-induced structure, diffraction peaks (10.0),

and (10.1

) ^ were used.

RESULTS

Analysis of the austenit- peaks, (111) and .90)

f

, showed a consistency in the relative intensi :es of each specimen; this implies little

variability of preferred orientation among the specimens.

The intensity traces of alloys 5 and 9, strained at 76 K to 0.05 and

0.10, are shown in Figure 1. The hep diffraction peaks us~d for this analysis were the (10.0)^ and (10.1)^. These hep peaks and most of the fee and

bcc peaks do not overlap; this enables accurate estimates of integrated intensity. In alloys 1, 3, and 5, the (110)^ peak, representative of straininduced bcc martensite, overlaps the (111)^ peak.

The relative intensities of the hep peaks were estimated by using both

computer-assisted curve fitting and graphical analyses. The curve fitting

to Lorentzian profiles used least-squares analyses and then integrated the

area within the final curve. From graphs such as that in Figure 1, peak

heights and half-widths were measured; their product was used as a relative

measure of integrated peak intensity. Both techniques produced very simila

data, which agreed within 10%.

The hep-martensite peak intensities were averaged. The average values

of these peak intensities are plotted versus atomic percentage of C and N

content ([C] + [N]) in Figure 2. In this plot, the fit of alloy-3 data is

not good. The relative peak intensities of all four alloys can be fitted

very well if the interstitial content is portrayed as 3[C] + [N]

,

which is

shown in the insert of Figure 2.

DISCUSSION

Figure 2 shows that the addition of both C and N to an Fe-20Cr-10Ni

austenitic alloy increases the fee stability with respect to formation of

strain-induced, hep martensite. If we consider the data of alloy 3, we st

conclude that C is more effective than N in austenite stabilization. Wh

the factor 3 [ C

]

+ [N] is used, the volume percent hep versus interstitial

content is clearly linear. Thus, C-and-N stabilization of austenite against

hep martensite is similar to their stabilization of austenite against bcc

66
Relative

Intensity

(Counts

per

second)

( 110)*'

Figure 1. Sample x-ray diffraction

peaks of fee ( 7 ), hep (e),

and bee (o’) structures

following 5 and 10% plastic deformation at 76 K.

Figure. 2. Relative integrated x-ray peak intensities of selected

hep diffraction lines following 76-K plastic deformation

(at strains indicated) versus [C + N]

.

67
martensite formation. Usually, however, C and N have been judged to contribute about equally in empirically developed calculations of martensite-start

temperatures on cooling and straining. 1 ’ 2

The effect of C on the stacking-fault energy of austenitic stainless

steels has been studied. 3 The contribution of C to the stacking-fault

energy was 26 mJ/m 2 per wt.%, a modest increase. The effects of additions

of C and N ( [ C ] + [N] <_ 0.364 wt.%, [N] <_ 0.264 wt.%) on the stacking-fault

energy of an Fe-18Cr-10Ni alloy were examined with x-rays.

4 Within data

scatter, the stacking-fault energy was almost independent of [C] and [N]

,

perhaps increasing slightly with increasing interstitial contents. With

transmission electron microscopy, values of 24 ± 3 and 23 ± 4 mJ/m 2 were

measured for the stacking-fault energy of an Fe-17Cr-12Ni-2Mo-0 .02C alloy

with [N] of 0.08 and 0.25 wt.%., respectively. 5 Weak-beam microscopy measurements showed that [N] in excess of 0.24 wt.% (to 0.52 wt.%) had little

effect on the stacking-fault energy of an Fe-19Cr-7Ni-8Mn-0 .03C alloy, but

at 0.21 wt.% N, the stacking-fault energy was considerably higher. 6

Thus, in austenitic steels, [C] increased the stacking-fault energy

slightly. The role of [N] is not quite clear: either the stacking-fault

energy is independent of [N]*’ 5 or, if the one datum point at 0.21 wt.% N

of Stoltz and VanderSande 6 is strongly considered, stacking-fault energy

decreases with increased [N]

.

The increase of stacking-fault energy as [C] increases is consistent

with the suppression of hep martensite as [C] increases. But the constancy

or reduction of stacking-fault energy, compared with the suppression of hep

martensite as [N] increases, deserves further discussion. The suppression

of hep as [N] increases suggests that the free-energy difference between fee

and hep structures increases with increasing [N] . Yet, the stacking-faultenergy data suggest that the free-energy difference between the fee and hep

structures remains constant or decreases as [N] increases.

Following Ishida 7 and earlier works**’ the stacking-fault energy (y)

of an alloy consists of bulk (Yg), segregation (Yg), and magnetic (ym

) contributions:

7

b

+ + Y

M

( 1 )

In Fe-18Cr-10Ni austenitic alloys, the magnetic state at room temperature is paramagnetic for the fee and hep structures and ferromagnetic for

the bcc structure. On cooling, these alloys remain in the paramagnetic

state until the Neel temperature, about 40 K, is reached. 10 Therefore,

for this discussion we neglect Y^.

The bulk contribution to the stacking-fault energy of alloys is represented as the sum of the difference of the bulk free-energy (G) and

interaction parameters (W) of the alloy components. If we assume that

interstitial N and the Fe-Cr-Ni austenite matrix form a two-component (A,N)

system that can be described by Ishida' s regular-solution model, then

68
~ -JILUUjjW

r B 2V 2/3 AG

f-h

(1 -V * a«TV <V-v ( 2 )

where X^ is the concentration of N, and V is the molar volume. The suppression of hep with increasing [N] (Fig. 2) implies an increase in the

free-energy difference between the two structures. Thus, with N addition,

the sum of the first two terms of Equation 2 is expected to increase. If

the change of the interaction parameter is positive (implying increased

order of N in the fee phase), the third term of Equation 2 also increases

*r^ with increasing [N] . When W is zero (in^tjie ideal-solution model), this

third term is inconsequential, but, when W^ is negative, the third term

of Equation 2^m^y balance the positive contribution of the first two terms.

A negative W implies that N in the hep structure is more ordered than in

the fee structure. 9

The fault-segregation contribution, T^, to the stacking-fault energy

depends on changes in the chemical-free energy (AG ), elastic energy (AG ),

and surface-free energy (AG ): s ’s'zk (3)

where

AG

c (1-XW ) In

N

(4)

h h and is the concentration of N in the fault. If f X^, AG

c decreases

only when W^ is negative, that is, an increased concentration of N species

is ordered in the fault.

The elastic-energy term is

dV'

,dXj

(5)

where u is the shear modulus, and v is Poisson's

addition there ig an attendant volume increase. 1 to increase if X^ # X^.

ratio. 7 With nitrogen

Thus, AG

g

is expected

To a first approximation, the surface free-energy term

7,9 is given by

AG

s - W

4 AN 4-h ( 6 )

and AG

g

is directly proportional to the interaction parameter if segregation

is present.

In summary, Equations 1 through 6 are used to describe interstitial

alloying effects on stacking-fault energy. It follows that the apparent

divergence of the phase transformation and stacking-fault-energy data with

69
increasing [N] may be explained qualitatively if N segregation with increased ordering at stacking faults occurs. Increased N ordering at stacking faults decreases the hep interaction parameter, providing negative terms

to a regular solution model of stacking-fault-energy dependence on structure

free energy, composition, and interaction parameters. Nitrogen ordering

could be achieved when N occupies either split-interstitial sites or preferentially couples with another alloying element, such as Cr. xl

REFERENCES

1. R.P. Reed, Alloying, p. 225, American Society for Metals, Metals Park,

Ohio (1988).

2. Handbook of Stainless Steels, ed. D. Peckner and I.M. Bernstein,

McGraw-Hill, New York (1978).

3. P.J. Brofman and G.S. Ansell, Metall. Trans. 9,879 (1978).

4. H.M. Ledbetter and M.W. Austin, Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures - VIII, p. 271, ed . R.P. Reed,

NBSIR 85-3025, National Institute of Standards and Technology, Boulder,

Colorado (1985).

5. R. Taillard and J. Foct, Proc. "HNS-88" (Lille, France), ed. J. Foct

and A. Hendry, Institute of Metals, London (1989).

6. R.E. Stoltz and J.B. VanderSande, Metall. Trans. 11,1033 (1980).

7. K. Ishida, Phys . Status Solidi 36a, 717 (1976).

8. J.P. Hirth, Metall. Trans. 1,2367 (1970).

9. T. Ericsson, Acta Metall. 14,1073 (1966).

10. E.R. Jones, T. Datta, C. Almasan, D. Edwards, and H.M. Ledbetter,

Mater. Sci. Eng. 91,181 (1987).

11. R.P. Reed, J. Met. 41,16 (1989).

70
LOW- TEMPERATURE PROPERTIES OF

HIGH-MANGANESE AUSTENITIC STEELS*

R. P. Reed, P. T. Purtscher, and L. A. Delgado

National Institute of Standards and Technology

Boulder, Colorado

ABSTRACT

The low- temperature properties of high-Mn austenitic steels

are reviewed. Emphasis is placed on strength, toughness, elastic

properties, thermal expansion, and magnetic properties, those

properties most critical for materials used in cryogenic

structures. Strength and toughness parameters and mechanisms are

discussed, including alloying, martensitic transformations, twinning, and grain size. Information is included from many recent

studies, particularly Japanese studies, of alloy development for

superconducting magnet applications at 4 K.

STRONG STRUCTURAL ALLOYS are required to withstand the large magnetic

forces of superconducting magnets. Since structural failure in these applications would result in unacceptable loss, these structural alloys must also

have excellent fracture toughness. Many of the low- temperature applications, such as those in magnetic-fusion-energy reactors and high-energyphysics magnets, require materials with high elastic properties and low

thermal and electrical conductivities. Austenitic steels have been almost

exclusively selected for these uses. Recently, the Japan Atomic Energy

Research Institute^) has stimulated the Japanese steel companies to develop

even stronger (1200-MPa yield strength at 4 K) and tougher (200-MPaVm fracture toughness at 4 K) austenitic steels for future use for structural constraint of the large magnetic fields. The impetus to achieve these goals in

Japan, coupled with research at the National Bureau of Standards and other

laboratories in the United States, has led to development of even better

Fe-Cr-Ni and Fe-Mn -Cr austenitic steels. The compositions of the high-Mn

steels that have been recently developed are summarized in Fig. 1. Those

with high Cr contents meet the Japanese corrosion-resistance requirements;

those with low Cr contents were developed prior to the corrosion-resistance

requirement or were developed for other applications. All contain

sufficient Mn to stabilize the austenitic structure against martensitic

transformation on cooling.

This paper summarizes our understanding of the behavior of these highMn austenitic steels at cryogenic temperatures. The most significant design

parameters—strength, toughness, and physical properties—are emphasized.

* All compositions in this paper are in wt.%, unless otherwise

specified.

71
HIGH -MANGANESE AUSTENITIC STEELS

Figure 1. Portion of the Fe-Mn-Cr ternary diagram with compositions of

recently developed high-Mn austenitic steels.

STRUCTURE

Additions of Cr to Fe—Mn alloys serve two purposes: (1) improvement of

corrosion resistance and (2) retardation of the intergranular failure mechanism of the Fe-Mn binary alloys (discussed later) . Additions of Cr have

little affect on the austenite stability of the Fe-Mn alloys, as illustrated

in Fig. 2( 2 ). Neither the thermally induced a (bcc) and c (hep) transformations nor the strain- induced a' and c' transformations are significantly

affected. Additions of Cr greater than approximately 15% result in hightemperature 6 ferrite or a phase (at high Mn contents). Alloying with

relatively small amounts of Ni, C, or N is needed to suppress these hightemperature phases that are deleterious to fracture. Also, from Fig. 2,

notice that at Mn contents greater than about 22%, no strain- induced a' is

detected, even at 4 K. The strain- induced e

' -martensitic structure is

suppressed above about 32% Mn, and the thermal e -martensitic phase is suppressed above about 28% Mn( 4 ).

Although Fig. 2 is taken primarily from Horiuchi et al.( 2 ), similar

studies were conducted earlier by Namekata and Higashi( s

) at temperatures

ranging from 4 to 273 K. They also studied the effect of 6%* Ni addition to

Fe—Mn—Cr alloys and found that neither € nor e

was detected at Mn contents

higher than 25%.

72
50,

4 K

Y « y'+ S

Y 8 + cr

rY + a,a 1 /« «'+a'

10 20 30 40

MANGANESE, wt%

50

Figure 2. Portion of the Fe-Mn—Cr ternary diagram illustrating alloy

dependence of phases formed during deformation or cooling to

4 K( 2

)

Additions of Ni and C to Fe-Mn alloys enhance austenite ( 7 ) stability.

The Fe-rich portion of the Fe-Mn—Ni ternary diagram( 5

) containing both

thermally and strain- induced phases at 76 K is shown in Fig. 3. From this

figure, it is apparent that the addition of about 10% Ni to alloys having

greater than about 15% Mn is sufficient to prevent the formation of both a'

and £ ' formation at low temperatures . Even less Ni is required for

austenitic thermal stability if C, N, or even Cr is present. Similarly, C

additions increase 7 stability( 6

) , as illustrated in Fig. 4. Although a

ternary Fe-Mn -N phase diagram was not discovered in review for this paper, N

additions are also expected to increase austenite stability.

Thus, to achieve austenite that is stable with respect to martensite,

Mn, Ni, C, or N is added. To prevent high- temperature phases, Ni, C, or N

is added and Cr is kept to about 15% or lower. If corrosion resistance is

desirable, 10 to 15% Cr must be added.

STRENGTH

Mechanisms

There are four distinct deformation mechanisms in high-Mn austenitic

steels at low temperatures: slip, deformation twinning, transformation to € martensite, and transformation to a' martensite.

73
50,

Figure 3.

Figure 4.

Portion of the Fe-Mn-Ni ternary diagram illustrating alloy

dependence of phases formed during deformation or cooling to

76 K( 6

)

.

Portion of the Fe-Mn- C ternary diagram illustrating alloy

dependence of phases formed during deformation or cooling to 300

K( 6 ).

74
SUs. - In very stable austenite or at higher temperatures, the only

deformation mechanism Is slip. Figure 5, which Is based on data from Remy

and Pineau( T

), shows that slip is found at high temperatures and high Mn

contents. Additions of Cr, Ni, C, and N reduce the temperature and the Mn

content at which slip is the dominant deformation mechanism. These alloying

elements increase austenite stability with respect to a' martens -ite formation and are expected to increase (or, at least, not to decrease) the

stacking- fault energy. Since slip occurs by dislocation nucleation and

motion, the factors that affect dislocation dynamics control the flow

strength in this regime. These factors include long-range interactions with

dislocation cells or grain boundaries; shorter range, temperature -dependent

interactions with interstitials such as C and N; and dynamic recovery.

Deformation Twinning — At lower temperatures or Mn contents , deformation

twinning is observed. Twins have been associated with bundles of stacking

faults aligned on (111}^ planes. Figure 6 shows the etched surface of an

Fe—22Mn—13Cr—5Ni— 0.22N alloy with 10% elongation at 76 K. The broad bands

are thought to represent deformation twinning, since no e' was detected

using x-ray diffraction, no o' was detected using magnetic measurements, and

simple slip bands would not be expected to etch. As shown in Fig. 5, the

addition of C to Fe-Mn alloys tends to increase the temperature range of

deformation twinning, but the addition of Cr decreases the range of twinning

while increasing the occurrence of «' martensite. Little is known of the

quantitative effect of deformation twinning on stress—strain characteristics, although the results of a study by Tomota et al.(*) of Fe—31Mn and

Fe-36Mn- alloy flow strengths imply that twinning does not substantially

alter the tensile behavior, as slip does.

Transformation to g' Martensite — Transformation to c' martensite, rather

than slip or twinning, occurs at lower temperatures and Mn contents

(Fig. 5). Formation of «' martensite during deformation increases the rate

of work hardening( #

) . The platelets of c ' on {111) 7 are thought to segment

the austenite grains, considerably obstructing the dislocation motion

through effective reduction in grain size. This effect of c' formation on

stress-strain behavior is illustrated in Fig. 7. In the absence of a' formation, strain- induced «' continuously increases with a linear dependence on

strain (Fig. 7)( 2 ).

Transformation to a' Martensite — At even lower temperatures or Mn contents

(see Figs. 2 through 4), a' martensitic transformation is strain induced.

In a manner similar to Fe-Cr—Ni alloy deformation behavior(®), a' formation

has two effects( 8 ), which are illustrated in Fig. 7. At strains usually

less than 0.10, a' forms at slip-band intersections as laths with a <110>7

axial direction corresponding to the common direction of {111}~ intersection. These laths effectively reduce the flow stress by relieving

stress, transmitting dislocation pile-ups, or both. Subsequent a' formation

acts in the opposite direction, increasing the rate of work hardening. At

larger strains, presumably the slip-band intersection sites for a' formation

have been exhausted and o' forms from c'. This produces a maximum in the

dependence curve of e' on strain. Thus at low temperatures, metastable

polycrystalline austenitic steels can have characteristic stress—strain

curves, very similar in shape to single -crystal curves, with three stages( 10 ).

75
Figure 5. Deformation mechanisms in Fe—Mn—(Cr or C) alloys as a function of

temperature and Mn content ( 7 )*.

Figure 6. Photomicrograph of Fe—22Mn—13Cr—5Ni— 0.22N austenitic steel

deformed in tension to 10% elongation at 76 K.
Figure 7. Stress—strain and associated martensite transformation characteristics for austenitic steels at low temperatures.

Alloying Effects — Although solid- solution alloying elements, such as Cr

and Ni, may affect the flow or fracture properties by influencing the

relative stability of the austenite (thus affecting o' , e', or twinning

deformation mechanisms) , they contribute virtually nothing to the

flow strength of stable high-Mn austenitic steels (whose deformation

mechanism is slip). This lack of influence is understandable, because these

elements have effective volumes close to Fe (Vjji/Vpe - 0.93, Vcr/Vpe -

1.02)

(

ll ). This similarity results in similar lattice parameter changes

with alloying additions. However, additions of the much smaller elements C

and N, which presumably occupy interstitial lattice sites, increase strength

77
considerably at low temperatures. The effective volume of C relative to

Fe is 0.34; that of N, 0.30( 11 ). In Fe—Cr—Ni alloys, the smaller N atom

produces larger changes in the lattice parameter (a): at 295 K, (1/a)

(da/dc) - 0.218 for C and 0.240 for NIC 11

), where c is composition in at.%.

We expect a similar lattice -parameter dependence in high-Mn austenites.

We conducted regression analyses on all high-Mn austenitic steel data

at 295, 76, and 4 K that we could gather( 12 ' 26

) . We defined high-Mn steel

as any steel having an austenitic structure on cooling to low temperature

and Mn as the principal alloying element.

The dependencies of yield strength (<7y, in MPa) on C and N content (

[ C

]

and [N] , in wt.%) and grain size (d, in were analyzed using the

expression

- o 0 + a[C] + b[N] + led"

1*, (1)

where o 0 is the strength before C and N additions.

The results are listed in Table 1 and plotted in Fig. 8. Brackets

drawn about the N and C coefficients in Fig. 8 represent one standard

deviation; the brackets about the data points for the ratios of N/C coefficients represent the estimated fractional error (F.E.), using

F.E. - ± (b 2 Pc

s + a*Pn2 )*Vab, (2)

' where Pc and Pn , the probable errors in a and b, were taken as one standard

deviation (Table 1).

Notable results are: (1) The temperature dependencies of the C and N

strengthening contributions are larger than that of a0 , implying shorter

range stress fields from the interstitials. (2) The ratio of these contributions are plotted versus temperature in Fig. 8 along with the ratios for

identically determined Fe—Cr-Ni alloy data( 27 * 28 ). It is apparent that the

relative contribution from N is higher in the Fe— Mn—Cr alloys than in the

Fe—Cr—Ni aus teni te s

.

Table 1. Regression Analysis of C, N, and Grain-Size Contributions

to Tensile Yield Strength in Eq. 1.

Number Coefficients Standard Deviation

Temperature of Data (MP,/«.%) (MPa/wt.%)

Points °o a b k °y *c a b k

4 58 551 773 2140 1140 77 52 142 178 161

77 25 264 315 1580 1060 76 64 44 145 454

295 66 118 132 683 462 53 23 24 46 89

78
Figure 8. Regression analyses results for the dependence of yield

strength on C and N content and grain size at low temperatures .

Furthermore, in the Fe—Cr-Ni alloys, N strengthens more than C at low

temperatures; in the Fe— Mn-Cr alloys the reverse is true. (3) Comparison

of the coefficients with results of previous Fe-Cr—Ni regression anal79
yses( 27 * 28

) indicates that the magnitude of the strengthening contribution

from C and N is less in Fe-Mn-Cr alloys than in Fe-Cr-Ni and Fe-Cr-Ni-Mo

(316- type) alloys. However, the smaller contribution to interstitial

strengthening in the Fe-Mn-Cr alloys is offset by two factors: the strength, a 0 , is almost twice that of Fe—Cr—Ni alloys (see Table 2) at 4 K, and

the solubility of C and N in Fe-Mn-Cr alloys is about twice that in Fe-Cr-Ni

alloys. Lula( 2 ®) lists the predicted and measured N solubilities at 1873 K

for Fe-15Mn-15Cr and Fe-15Ni-15Cr alloys as 0.29 and 0.13 wt.%, respectively. Both Mn and Cr additions increase the solubility of N in austenite;

Ni addition results in slightly lower N solubility.

The addition of N to Fe-Mn-Cr austenites increases their strength by

about a factor of three more than the addition of C (Fig. 8). This very

significant distinction between C and N strengthening is not predictable on

the basis of size or modulus misfit(®), which leads to the suspicion that N

is positioned differently in the lattice than C, possibly in pairs (®) or

ordered within short range

(

30 ).

GRAIN SIZE DEPENDENCE

The published data( 1

5

" 1

7

* 20 * 25 * 3

1

" 33

) on the effects of grain size on

tensile yield strength, using the Hall— Petch relationship, are included in

Fig. 9. Here,

°y “ a o' + kd* 15

, (3)

where ct 0

' includes all grain- size -independent contributions.

Table 2. Nitrogen and Carbon Contributes to Tensile Yield Strength.

Source Alloy 4 K

Nitrogen

76 K

(MPa/wt . %)

295 K

Carbon

4 K 76 K

(MPa/wt . %)

295 K

This paper Fe—Mn—Cr 2140 1580 680 770 320 130

Simon, Reed,

1986( 28

) Fe-Ni-Cr 3190 1660

Reed, Simon,

1984( 27 )*

Fe-Ni-Cr 3400 2080 420 1700 840 830

Simon, Reed,

1986( 23

) Fe-Ni-Cr-Mo 2780 1690

* Did not include a grain-size term in the analysis.

80
The results of our regression analyses (using Eq. 1) in calculating the

grain-size dependence, k, are also included in Fig. 9. A typical 4-K value

of k for Fe-Cr-tfi alloys

(

2S

) is 1090 MPa/jim'

1*; for Fe—Cr—Ni-Mo alloys, 810

MPa/zim'^C 34) . The lower average values for these alloys are outside the

data spread of individual Fe—Mn-Cr- alloy studies. The increased grainboundary contribution of the Fe-Mn—Cr alloys implies increased alloy segregation at their grain boundaries, presumably by C and N.

The dependence of flow strength on grain size is quite strongly dependent on temperature, increasing by a factor of 2 from 295 to 4 K. Purer

face-centered cubic metals, such as 99.95% Cu, exhibit about a 30% increase

in k from 295 to 4 K( 36 ). Thus, the stronger temperature dependence of the

grain size influence on flow strength also probably reflects more impurity

segregation at grain boundaries of the austenite structure. However, as

illustrated in Fig. 9, there is no clear trend of k as a function of

C and N content.

Figure 9. Summary of literature data on the Hall- Petch constants of highMn austenitic steels at low temperatures. Also included are the

k results of the regression analyses of this study.

81
TEMPERATURE DEPENDENCE

From analyses of the temperature dependence of many austenitic steels,

we found that for most alloys at low temperatures , the log of the yield

strength is linearly and inversely related to the temperature. This is

illustrated in Fig. 10 for several high-Mn alloys( 22

), and can be expressed

as

Oy - a o "exp (—AT) , (4)

where ct 0

" is the yield strength at absolute zero and A is a constant, the

slope of In ay versus temperature, T. This slope is proportional to the

estimated stacking- fault energy of the austenitic alloys( 22 ).

Near room temperature, most alloys deviate from logarithmic dependence

(Fig. 10), but there is no other evidence that this deviation signifies a

change of the deformation mechanism. Many of these alloys are completely

stable: there is no strain- induced martensite. The macroscopic contours of

the stress—strain curves remain the same. Obviously, fundamental studies

are needed to assess the significance of the temperature -dependence trends.

TOUGHNESS

MECHANISMS — The toughness of binary Fe-Mn alloys is degraded by intergranular fracture, even in stable austenite( 4 » 1S

) . Recent studies( 4 » 1S

) have not succeeded in detecting segregation at the grain boundaries; thus,

at present, there is no remedy for this fracture mechanism in the binary

alloys. However, additions of the order of 5% Cr or Ni or small amounts of

C do inhibit intergranular fracture.

Figure 10. Temperature dependence of austenitic alloy yield strengths at

low temperatures

.

82
At low temperatures, Fe-Mn-Cr and Fe-Mn-Cr— Ni alloys with high N contents or low austenite stabilities exhibit faceting of the stacking -fault

clusters on {111 } -y in impact or fracture toughness tests

(

4 * 36 " 3 ® j . This

brittle -looking faceting is similar to that reported by Tobler and Meyn( 39

) as slip-band cracking on compact- tension specimen fracture surfaces in highN alloys. These facets usually have other {111}^ traces on their surfaces

(Fig. 11a). The fracture surfaces following fracture toughness tests at low

temperatures are seldom dominated by these {111}^ facets; elongated, shallow

dimples are the prevalent morphological structure (Fig. 11a). higher

temperatures or in more stable structures, the faceting is reduced and the

elongated dimples are dominant on the fracture surfaces (Fig. 1 lip ) . This

change is usually associated with increased toughness. Perhaps the reduction of fracture -surface faceting is assisted by Ni addition. This is the

case for Fe—Cr—Ni alloys

(

40 ). Yamamoto et al.( 20

) and Suemune et al.( 41) have also found that Ni increases the toughness of the Fe—Mn—Cr alloys, but

the fracture mechanisms in their studies were not identified. It is still

uncertain whether the role of Ni in Fe—Mn—Cr austenite fracture is to increase toughness by increasing the stacking- fault energy (suppressing e'

formation, deformation twinning, and stacking -fault clusters) or to affect

the toughness of alloys that fail by void growth and coalescence.

At higher temperatures, decreased strengths, higher Ni contents, or

lower stress concentrations, typical ductile dimple fracture is prevalent

(Fig. 11c). Nonmetallic inclusions are associated with most ductile dimples. Yamamoto et al.( 20

) and Suemune et al.( 41

) have both obtained excellent correlations between reduced S content (to reduce MnS inclusions) and

toughness at low temperatures for Fe-22Mn— (13—15)Cr— (l-7)Ni austenitic

steels. Apparently, the conditions, such as those illustrated in Fig. lib,

reflect the fracture morphology transition between faceting and dimpling.

In this transition, many of the shallow, elongated dimples do not appear to

be nucleated at inclusions.

Strength Dependence

The dependence of toughness on strength in high-Mn austenitic alloys is

similar to that in Fe-Cr—Ni alloys. The inverse dependence of yield

strength of a series of Fe—20Cr—lONi—C+N alloys (the shaded band) is shown

in Fig. 12; the lines on each side of the band indicate the data spread of

all Fe—Cr—Ni base metals that we have measured at 4 K. The Kjc (J) fracture

toughness values plotted in this figure were converted from J- integral

measurements using

[KIc (J)] 2 - EJ Ic , (5)

where E is Young's modulus. The compact tensile specimens measured were

oriented in the TL direction (Jjc measurements are about 10 to 20% lower in

this orientation than in the LT orientation) . The tensile yield-strength

data in Fig. 12 were measured in the transverse direction. The individual

data points portray our measurements of high-Mn austenites. The several

data points of the solution- treated alloys 18Mn—16Cr—5Ni—0 . 2N and 22Mn—13Cr5Ni—0.2N reflect test and material variability. The location of these highMn- alloy data points on the high side of the trend lines probably reflects

increased steel cleanliness. The fracture appearance of these high-Mn

alloys varies: high-yield-strength alloys have more {111}^ facets; lowyield-strength alloys have larger dimples.

83
Figure 11. (a) Fracture surface of a compact- tension specimen of

Fe—8Mn—20Cr—7Ni—0 . 37N alloy, tested at 4 K, displaying

{ 1 1 1 } -y brittle facet and elongated, shallow microvoid

dimples

.

(b) Fracture surface of a compact- tension specimen of

Fe—22Mn—13Cr—5Ni alloy tested at 4 K.

(c) Fracture surface of tensile specimen of Fe—22Mn—13Cr—5Ni

alloy tested at 4 K.

84
Figure 12. Dependence of fracture toughness on tensile yield strength at

4 K for (1) a series of Fe—Cr—Ni alloys with controlled amounts

of C and N (shaded band), (2) all Fe—Cr—Ni austenitic steels

measured at NBS (wide band) , and (3) high-Mn austenitic steels

measured at NBS

PHYSICAL PROPERTIES

At low temperatures, high-Mn austenites are antiferromagnetic; at high

temperatures, they are usually paramagnetic. The magnetic transformation at

their Neel temperature significantly influences their elastic properties,

thermal expansion, and magnetic susceptibility.

Magnetic susceptibility is probably the best indicator of the Neel

temperature. At the transition temperature, magnetic susceptibility peaks

sharply (Fig. 13). The data of Collings( 42

) show dependence of the Neel

temperature on composition: The Neel temperature of 13Mn—18Cr-3Ni is 160 K

higher than that of 5Mn—22Cr-13Ni

The transition to an antiferromagnetic state produces a softening of

the elastic properties. This softening is relatively large, as illustrated

by Ledbetter's data( 43

) in Fig. 14 for a Fe-26Mn—5Cr alloy. Both Young's

modulus (E) and the shear modulus (G) decrease about 8% through the Neel

temperature. The bulk modulus (B) is less affected; Poisson's ratio (i/)

increases about 8%.

85
(a) (b) (c)

Figure 13. Magnetic susceptibility dependence on temperature and alloy

content for three Fe—Mn— Cr—Ni alloys [Collings( 38

)

] . The

three susceptibility scales refer to alloy designations

(a) Fe—22Cr—13Ni— 5Mn, (b) Fe-21Cr-6Ni-9Mn, and

(c) Fe- 18Cr—3Ni—13Mn.

Figure 14. Elastic properties for an Fe—26Mn—5Cr alloy at low temperatures

[Ledbetter

(

3 ®

)

]

.

86
The magnetic transition also strongly affects thermal expansion.

Below the Neel temperature (Tjg) , there is positive magnetostriction in the

antiferromagnetic state, which reduces the thermal expansion coefficient.

Even at temperatures up to 100 K above the transition temperature, there is

limited magnetic ordering that lowers the coefficient.

Richter and Pepperhoff

(

4

) have measured the thermal expansion

characteristics of a number of Fe—Mn base alloys at low temperatures. They

observed that C and Cr additions tend to lower Tjj and reduce the volume

magnetostriction (Fig. 15). In the Fe—Mn binary system, Mn additions of 32

to 50% increase Tjj, and Mn additions of 32 to 40% increase the volume magnetostriction; further additions of Mn decrease both terms. Thus, the

amount of thermal expansion depends critically on the location of the Neel

temperature. In general, the lower the Neel temperature, the higher the

expansion coefficient. Sasaki et al.( 46

) have developed an empirical relation suggesting a nearly linear, inverse dependence of the thermal expansion

coefficient on Tjg.

Figure 15. Effects of alloying on Neel temperature and volume magnetostriction of Fe-Mn-base alloys( 40 ).

87
In principle, high-Mn austenitic alloys may be designed to produce

varying contraction characteristics at low temperatures. If an alloy is

desired with a lower thermal contraction matching that of a low-alloy

ferritic steel, the Mn content should be high enough to ensure that Tjj is

above room temperature (Fig. 16). If a high-Mn steel is to approximate the

contraction characteristics of Cu, Al, or Fe—Cr—Ni alloys, the Mn content

should be lower.

SUMMARY

The strength, toughness, and selected physical properties of high-Mn

austenitic steels at low temperatures have been reviewed. Compared with

Fe—Cr—Ni austenitic steels, the high-Mn steels have several distinct characteristics .

The solubility of N and C is higher, which enables the attainment of

higher strengths. More hep martensite is induced by deformation.

The hep phase increases the rate of work hardening during deformation. In

toughness tests, increased fracture faceting on {111}^ is observed in lower

stacking- fault-energy alloys. Binary Fe-Mn alloys are more susceptible to

intergranular fracture. Therefore, high-Mn austenites are easier to

strengthen than Fe—Cr—Ni austenites, but more troublesome to toughen.

The elastic properties and thermal expansion are strongly influenced

by the magnetic transition from the high- temperature paramagnetic state to

antiferromagnetism at low temperatures.

Figure 16. Percent contraction of various face-centered cubic structural

alloys on cooling.

88
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,

Kawasaki, Japan (1985)

20. Yamamoto, S., N. Yamagami, and C. Ouchi, in "Advances in Cryogenic

Engineering — Materials," vol. 32, pp. 57—64, Plenum Press, New York

(1986)

89
21. Tone, S., M. Hiromatsu, J. Numata,

T. Horiuchi, H. Nakajima, and S. Shimamoto, in "Advances in Cryogenic Engineering - Materials," vol. 32, pp. 89- 96, Plenum Press,

New York (1986)

.

22. Reed, R.P. and J.M. Arvidson, in "Advances in Cryogenic Engineering -

Materials," vol. 30, pp. 263—70, Plenum Press, New York (1984).

23. Kaji, H. , M. Hiromatsu, S. Tone, M. Shimada, and S. Shimamoto, "The

Development of a 22Mn-13Cr-5Ni Nonmagnetic Steel Plate for Cryogenic

Applications," Kobelco Technical Bulletin 1093, Kobe Steel Co., Kobe,

Japan (Nov. 1985).

24. Suemune, K. , T. Sakamoto, T. Ogawa, T. Okazaki, S. Maehara,

H. Nakajima, and S. Shimamoto, in "Advances in Cryogenic

Engineering - Materials," vol. 34, pp. 123-29, Plenum Press, New York

(1988) . 25. Ouchi, I., Nippon Kokan K.K. , Kawasaki, Japan; unpublished data

(1988).

26. Muira, R. , K. Ohnishi, H. Nakajima, Y. Takahashi, and K. Yoshida, in

"Austenitic Steels at Low Temperatures," pp. 287—93, Plenum Press,

New York (1983)

.

27. Reed, R.P. and N.J. Simon, in "Advances in Cryogenic Engineering —

Materials," vol. 30, pp. 127—36, Plenum Press, New York (1984).

28. Simon, N.J., and R.P. Reed, J. Nucl. Mater. 141—143 . 44-8 (1986).

29. Lula, R.A. , in "Manganese Stainless Steels," pp. 30—1, The Manganese

Centre, Paris, France (1986).

30. Bynes, M.L.G. , M. Grujicic, and W.S. Owen, Acta Metall. 35, 1853—62

(1987).

31. Kato, T. , M. Fujikura, S. Yahagi, and K. Ishida, Tetsu- to-Hagane 67,

587-95 (1981).

32. Yoshimura, H. , H. Masumoto, and T. Inoue, in "Advances in Cryogenic

Engineering — Materials," vol. 28, pp. 115—25, Plenum Press, New York

(1982).

33. Ogawa, R. and J.W. Morris, in "Advances in Cryogenic Engineering -

Materials," vol. 30, pp. 177—84, Plenum Press, New York (1984).

34. Reed, R.P., N.J. Simon, P.T. Purtscher, and R.L. Tobler, in

"Proceedings, International Cryogenic Engineering Conference,"

pp. 786— 90, Berlin, Butterworths , London, England (1986).

35. Reed, R.P. , R.P. Walsh, and F.R. Fickett, in "Advances in Cryogenic

Engineering - Materials," vol. 34, pp. 299-308, Plenum Press, New

York (1988).

36. Thompson, L.D. , in "Mechanical Properties and Phase Transformations

in Engineering Materials," pp. 391-408, AIME Metallurgical Society,

Warrendale , Pennsylvania (1986).

37. Kato, T. , S. Fukai, M. Fujikura, and K. Ishida, Trans. ISIJ 16,

673-79 (1976).

38. Defilippi, D.P., K.G. Brickner, and E.M. Gilbert, Trans. AIME 245,

2141-48 (1969).

39. Tobler, R.L. and D. Meyn, Cleavage- like fracture along slip planes

in Fe—18Cr—3Ni—13Mn austenitic stainless steel at liquid helium

temperature, Metall. Trans.; accepted for publication.

40. Purtscher, P.T., R.P. Walsh, and R.P. Reed, in "Advances in Cryogenic

Engineering — Materials," vol. 34, pp. 379—386, Plenum Press, New

York (1988).

90
41 . Suemune, K. , K. Sugino, H. Masumoto, H. Nakajima, and S. Shimamoto,

in "Advances in Cryogenic Engineering - Materials," vol. 32,

pp. 51—6, Plenum Press, New York (1986).

42. Collings, E.W. , in "Advances in Cryogenic Engineering — Materials,"

vol. 26, pp. 37—47, Plenum Press, New York (1979).

43. Ledbetter, H.M. , Phvslca 119B . 115-8 (1983).

44. Richter, F. and W. Pepperhoff, Arch. Eisenhuttenw. 47, 45-50 (1976)

45. Sasaki, T. , K. Watanabe, K. Nohara, Y. Ono, N. Kondo, and S. Sato,

Trans. Iron Steel Inst. Jap. 22 . 1010—20 (1982).

91
.

-

.

mm
DISCONTINUOUS YIELDING IN AUSTENITIC STEELS

AT LOW TEMPERATURES

R.P. Reed and N.J. Simon

National Institute of Standards and Technology

Boulder, Colorado,

The effects of strain rate and coolant state on the discontinuous yielding characteristics of a number of austenitic steels

at 4 K were studied. In liquid helium at 4 K, a transition in

initiation stress and strain as a function of strain rate was

identified. In gaseous helium at 4 K, discontinuous yielding

began at significantly lower stresses and strains than in liquid

helium. Results are interpreted in terms of macroscopic heat

balances and localized generation of heat from moving dislocations .

INTRODUCTION

At very low temperatures, the specific heat and the thermal conductivity of stainless steels and many other alloys are very low. Thus, it

is difficult to maintain the low temperature of the alloy when a source of

heat is applied. During plastic deformation in liquid helium (4 K) , only a small amount of the applied energy is stored in the alloy as defect

formation energy; most is converted to heat, which must be removed. Reed

and Walsh 1 described two distinct modes of transient heating that occurred

as a function of strain rate during straining of austenitic steels at 4 K:

thermal waves and thermal spikes. Thermal waves were associated with

macroscopic (overall) warming of the specimen; thermal spikes were

associated with discontinuous yielding and microscopic (local) specimen

heating. They concluded that (1) relatively little (-10%) energy is stored;

(2) much more heat is transferred from the deforming tensile specimen to the

liquid than is conducted through the specimen ends; and (3) at high strain

rates, the transition from nucleate to film-boiling heat transfer at the

liquid interface causes thermal waves, an indication of overall specimen

heating.

Thermal spikes as high as 60 K have been reported. 1-6 Associated with

discontinuous (serrated) yielding, they occur at low strain rates. In tests

in which thermal spikes are prevalent, the temperature of the deforming

specimen remains at 4 K prior to the strain instability; no precursor temperature increases have been detected. We think discontinuous yielding may

be a two-step process: (1) the triggering event that probably occurs at the

microscopic level and (2) the subsequent dislocation avalanche at very high

93
strain rates at the macroscopic level. This subsequent plastic deformation

occurs in bands across the specimen, similar to Luders bands. The bands are

about 10” 2 m wide with a typical extension of 1.3 x 10” 4 m. The local temperature rise, confined to the strained region, is adequately described by

assuming adiabatic conditions.

Previous studies 1 ” 6 of discontinuous yielding in austenitic stainless

steels described temperature rises of about 50 K in the deformation region

with very high local strain rates. Attempts have been made to analyze and

characterize discontinuous yielding at very low temperatures. Originally,

Basinski 6 reported low- temperature discontinuous yielding in aluminum alloys

and suggested that localized temperature rises, due to adiabatic conditions,

caused unstable plastic deformation. Later, on the basis of more macroscopic temperature measurements (maximum of 55 K during serration) , he

argued that thermal instability, rather than mechanical instability, was the

basis for the nucleation event. 7 This fundamental question has not yet been

resolved, because nucleation events probably have microscopic origins. Like

Basinski, Medvedev8 suggested a thermal heat balance between work of deformation and conduction from the specimen to account for the discontinuous

yielding process. Using computer simulation, Kubin et al

.

9 described the

low- temperature mechanical instabilities in terms of heat transfer to the

thermal bath and an Arrhenius rate equation. Shibata et al. 10 also used

computer simulation of unstable low- temperature deformation and found a balance between heat and mechanical instabilities similar to that described by

Kubin et al . , although a more realistic model for the stress dependence of

the exponential term was assumed. Like Reed and Walsh, 1 they considered

three elements in the heat balance: stored energy, heat conduction along the

specimen, and heat conduction to the coolant. Their results indicate that

a triggering event is unnecessary—the dynamics between mechanical work and

the thermal balance account for mechanical-thermal instabilities.

Moving dislocations cause local heating of alloys. Eshelby and Pratt 11

derived an expression (based on an internal source of dislocations) for the

local temperature rise at a distance, r^, from the dislocation in a metal:

where b is the Burgers vector; a, the applied stress; E, the average

dislocation velocity; k , the thermal diffusivity [thermal conductivity (A)

divided by the specific heat per unit volume of the metal]. For a number of

dislocations, n, spaced over a distance, i, on a glide plane at low

temperatures, AT is estimated to be

Armstrong et al. 12 pointed out that more localized slip-band formation, promoting fewer slip bands per active dislocation, results in larger temperature rises in Eq. 2. Nonuniform slip, confined to slip bands, is a trait of

face-centered cubic metals at low temperatures. The degree of slip

localization is thought to be affected by the stacking- fault energy.

94
The local heating of dislocation ensembles was studied by Malygin. 13

For absorption of heat along the moving dislocations, he considered only the

local heat capacity, not the local thermal conductivity. Using a Green's

function, he calculated individual temperature rises of 10" 2 K for moving

dislocations, a value less than Eshelby and Pratt's estimate from Eq. 2.

For an active Frank—Read source, Malygin calculates that the amount of heat

generated is Q - Q^n/3, where is the heat generated by the leading loop

and n is the number of loops. For a slip band with active sources on close,

parallel slip planes, local temperature rises were estimated:

AT - Wt'/Cv , (3)

where t' is the time, Cy is specific heat per unit volume, and W - cr^k ; - BT5/b and € - bpmD; pm represents the mobile dislocation density, and B,

a viscous drag stress. Temperature rises of 5 to 50 K can be calculated for

low specific heat and high dislocation velocities.

In this paper, we use the results of more than one hundred tensile

tests of austenitic stainless steels at 4 K to characterize stress and

strain for the initiation of discontinuous yields. For AISI alloy types

304L, 304HN, 310, and 316LN in liquid helium and types 310 and 304HN in a

gaseous helium environment, initiation stress and strain for discontinuous

yielding were measured at various strain rates. Alloy 304L has low strength

and low austenite stability; under plastic deformation, it transforms easily

to martensite. Alloys 304HN and 316LN are nitrogen- strengthened highstrength alloys; under large plastic deformations, they transform slightly

to martensite. Alloy 310 is completely stable. We report our measurements

in the form of heat-balance calculations.

EXPERIMENTAL PROCEDURES

Specimens of alloys 304L and 310 were machined from 1 . 9-cm-diameter bar

stock; specimens of alloy 304HN, from a 16-cm-diameter billet; specimens of

alloy 316LN, from 5.0-cm-thick plate. Alloys 304L, 304HN, and 316LN were

annealed; alloy 310 was tested in the partially sensitized condition.

Actual chemical compositions for these alloys were given previously; 1 nominally, in wt.% they are: alloy 304L, 18Cr-10Ni—lMn—0 . 02C—0 . 03N ; alloy

304HN, 20Cr-8Ni— 2Mn-0.06C-0.28N; alloy 310, 25Cr-21Ni-2Mn-0.09C-0.03N;

alloy 316LN, 18Cr- lONi-IMn—0.02C—0. 16N. The heat- treatment schedules,

hardness, grain size, 1 and microstructures 14 of the alloys were also

discussed previously. Round, tensile specimens 14 were machined to a 6. 35 -mm

diameter along a gage length of 41.9 mm, increasing with a radius of 9.52 mm

to 12 . 7 -mm -diameter threaded ends. The entire specimen length was 71.4 mm.

The cryostat, strain-gage extensometers , and temperature -sensing equipment have been described elsewhere. 14 * 15 Most tests were conducted in

boiling liquid helium; a few tests were conducted in gaseous helium (at 4 K)

with liquid-helium reservoirs above and below the specimen to regulate the

temperature. Elongation was measured over a specimen gage length of 3.8 cm.

Crosshead rates were varied from 0.001 to 2 cm/min. Reported strain rates

(from 4 x 10“ 6 to 4 x 10“ 3 s" 1 ) refer to the nominal rate of plastic

deformation within the reduced section of the specimens.

95
Figure 1 . The stress and strain at initiation of discontinuous yielding

and the tensile yield strength versus strain rate at 4 K for

alloy 304L.

EXPERIMENTAL RESULTS

For all alloys tested, the stress—strain curves at 4 K were initially

smooth and, subsequently, discontinuous yielding initiated. The effects

of strain rate on the stress, a^, and strain, to initiate the first

observable discontinuous yield in boiling liquid helium are shown in Figs. 1 through 3 for alloys 304L, 310, and 316LN, respectively. Clearly, at low

and high strain rates, is essentially independent of strain rate; at

intermediate strain rates, a transition occurs. All three alloys behave

similarly. Alloy 310 was tested in gaseous helium, as well as liquid

helium, at 4 K. Both a

and were considerably less in the gaseous environment (Fig. 2); in fact, the values are comparable to the high strainrate values that would be obtained in liquid helium. Another alloy, 304HN,

was also tested at 4 K in both liquid and gaseous helium at a low strain

rate (8 x 10“ 5 s" 1 ). The results are similar: in liquid, a

- 1730 MPa; in

gas, - 1440 MPa. The yield strength values, Oy (0.2% offset), were not

affected by the state of the cryogen.

96
Strain Rate (£), s

Figure 2. The stress and strain at initiation of discontinuous yielding

and the tensile yield strength versus strain rate at 4 K in

liquid helium (L^e ) for alloy 310. Also included are

measurements from testing in gaseous helium (Gjje ) at 4 K.

The absolute values of a £ for the same k range vary appreciably among

alloys. However, a

does correlate with Oy. Figure 4 shows versus Oy

for more than eighty measurements of austenitic steels in liquid helium.

Some steels are Fe—Cr-Ni-base alloys with varying carbon and nitrogen

content; some are high -manganese austenitic steels, usually with chromium

and nickel added; and many are experimental alloys from previous studies at

the National Bureau of Standards. The o^—Oy correlation is quite good; the

solid line depicts the relation (in MPa)

— 320 + Oy. (4)

All values were measured at low k within the high-o^, low-e plateaus illustrated in Figs. 1 through 3. The dashed lines were constructed from the

high- k -plateau measurements shown in these figures. In essence, increased

strain rates reduce the constant (to 110 from 320 MPa) in Fig. 4 and Eq. 4

but do not significantly change the linear dependence of a

on Oy.

97
0.10 1 1 ‘ 1 1 nil I 'll I I lllj. I I I I I I III I ^

0.08

c

j? 0-06

05

I

a04

:•= o.o2

c

0

-

M x

n

316LN, 4 K

X \x

\

\ V

\

V

*

1300 -

1200 * £ 1100

wf 1000

CO

<D E 900

•

800

%

700 ! 1

’ 1

X X

w X

\

X

a, Vs <

x\

\

/

Strain Rate (£), s

Figure 3. The stress and strain at initiation of discontinuous yielding

and the tensile yield strength versus strain rate at 4 K for

alloy 316LN.

DISCUSSION

Implications of the Experimental Results

From the experimental results, we consider the following points to be

characteristic of discontinuous yielding:

1. Both a

i

and differ radically in alloys 310 (Fig. 2) and 304HN when

measured in gaseous and liquid environments, even though the strain

rate and nominal test temperature were equivalent. This strongly

implies that surface heat transfer plays a major role in the initiation

of discontinuous yielding.

2. The transition from high a

(and e^) at low k to low (and c^) at

high € may be interpreted in two ways: (1) The high k effectively

increases the dislocation velocity:

k - bOpm , (5)

98
Figure 4. The stress at initiation of discontinuous yielding versus yield

strength for many austenitic steels at 4 K in liquid helium at

c • < 6 x 10-s s _1

. Dashed line is an estimate of dependency

of Of on a

y

for e > 5 x 10” 4 from the data for alloys 304L, 310,

and 316LN shown in Figs. 1 through 3.

and thus increases the internal dislocation heating. (2) The high k increases the rate of heat generated from the energy of deformation:

Edef " ( 6 )

where V is the volume of the deformed region and E^ef - *Edef»

following the nomenclature used by Reed and Walsh. 1 This increased

heating rate may result in a different mode of boiling heat transfer

(film rather than nucleate) to the liquid helium in a manner like that

suggested for the occurrence of thermal waves at high k . 1

3. The dependence of a ^ on Oy (Fig. 4, Eq. 4) is very interesting.

Consider that the flow strength, a, is the sum of contributions from

strain hardening, o^, and solution hardening, of.

O - C7 d

+ Of. (7)

The solution hardening contribution, o f, may be equated to the yield

strength, Oy. At the initiation of the first discontinuous yield,

o equals Of. Thus, o^ has a constant value of 320 MPa (from Eqs . 4

and 7). This constancy of o^, which represents dislocationdislocation interactions, suggests that, for initiation of

99
discontinuous yielding, a constant number of dislocations must be

accumulated. The usual representation of the effect of the total

dislocation density, (created by plastic deformation), on stress is

ad - aGbp*, (8)

where a is a constant, G is the shear modulus, and b is the Burgers

vector.

These implications require more careful consideration, but first, to assist

in understanding the discussion following, we briefly discuss heat transfer

from the specimen to gaseous and liquid helium.

Surface Heat-Transfer Mechanisms at 4 K

Two different heat- transfer regimes encountered in steady- state pool

boiling of helium at atmospheric pressure are indicated in Figure 5. This

figure presents typical boiling heat- transfer data abstracted from two

review papers. 16 ’ 17 It also shows a typical curve for heat transfer to

gaseous helium at 4 K, using the Raithby and Hollands's correlation for

laminar and turbulent flow at high Rayleigh numbers. 18 Heat-flux rate, Q/A,

is plotted against AT, the temperature difference between the solid surface

and the liquid. In the nucleate -boiling region, vapor bubbles are created

at nucleation sites on the solid surface. In the film-boiling region,

bubbles form so rapidly that a vapor blanket is established that prevents

liquid from contacting the surface. The heat- transfer coefficient,

h (- Q/A -AT) , is therefore lower than in the nucleate region. When the peak

nucleate -boiling heat- flux rate, (Q/A) max , is reached, the difference in

temperature, AT, between the specimen and the liquid may suddenly increase

by an order of magnitude.

Figure 5. Typical nucleate -boiling and film-boiling heat transfer in

liquid helium at 4 K (solid line) . Predicted heat transfer in

gaseous helium at 4 K (dashed line)

.

100
The value of (Q/A) max is typically about 10 4 W/m2

; it is nearly independent of AT and various geometrical and surface factors , which do change

the shape and position of the nucleate-boiling curve. For example, smoother

surfaces with smaller asperities shift the nucleate boiling curve in Fig. 5

to the right, raising AT by an order of magnitude without changing

(Q/A) max . 17 The slope, d log(Q/A)/d logAT can also increase. 19 However, a

change from a horizontal to vertical heat- transfer surface may increase

(Q/A) max only about 20%. 16,19 By reducing the size of miniature heaters in

liquid nitrogen, a decrease in (Q/A) max was attained; for normal heaters,

size was not a factor. 19 Increases in (Q/A) max to 10 s W/m2 or more have

been observed only with pulsed heat inputs, 20 that is, transient rather than

steady- state heat- transfer conditions.

Surface Heat Transfer

The influence of the test environment on the low- temperature discontinuous-yielding process implies that the energy—heat balance is not simply

confined to heat generation from moving dislocations in localized bands.

The deformation energy, E^ef. is either stored as energy, Egg. or dissipated

as heat, and the dominant heat- conduction path is through the specimen surface to the test environment

1 » 9 » 1

0

Thus, since energy is conserved at any

instant of time, 1

Edef “ eSE + Qs> ( 9 )

where Qs is the heat-dissipation and Ese “ Egg/6

Using a macroscopic approach, we can calculate E^ef as a i^i^* where V

is the deformed specimen volume to be considered. 1 From the values for

alloy 310 (Fig. 2), for example when q is 8 x 10“ 5 s" 1

, a

equals 1100 MPa

and a possible deformed volume is 7rr 2 t, where r is the specimen radius and t

is the "activated" thickness of the specimen cross section to be considered.

For comparison with Fig. 5, E^gf/Ag needs to be calculated (As - 27rrt is the

surface area of heat transfer). For these data, E^gf/Ag - 221 W/m2

, a rate

well below the critical value, (Q/A) max , of about 10 4 W/m2 for the

transition from nucleate -boiling to film-boiling heat transfer (Fig. 5). A

local warming of about 0.2 K is predicted by Fig. 5 for these conditions

when the specimen is tested in liquid helium. A heat transfer rate of 167

W/m2 is calculated for the gaseous -helium condition (from Fig. 2, q * 8 x

10” 6 s" 1

, a ^ - 830 MPa); Fig. 5 shows that this corresponds to a local

temperature rise of about 4 K. These diverse results suggest that a

different approach is needed to rationalize the measurements in liquid and

gas.

Consider the heat- transfer process across a portion of one grain (e.g.,

a slip band) at the specimen surface. Microscopically, the heat per unit

time and length from a moving dislocation is bcrU. 11 We consider an ensemble

of n moving dislocations (within a slip band); then the heat dissipation, Q,

from this ensemble is simply nbaUi^ (i<j - average dislocation length) . The

number of moving dislocations within the local area (A^g) of concern may be

conservatively estimated from the average mobile dislocation density, pm , as

n " Pm' Ai- Here,

101
“ P t Pn* ( 10 ) with pn the number of dislocations that have been generated but probably

have become immobilized in networks. The applied stress minus the stress at

which dislocations begin to generate (taken here as the yield stress ) is

,

from Eqs . 7 and 8

,

o — Oy — aGbp^. (11)

When Eqs. 10 and 11 are combined and a (- 1) is dropped,

Pm ~ (Gt>) 2 - pa

. (12)

Our transmission-electron-microscopy studies of austenitic stainless

steels with plastic-deformation strains less than about 0.10 found little

dislocation-network formation. Therefore, we assume that pm » pn and pn

can be eliminated from Eq. 12. Then,

Q

g

d

<7 i pi

d

A

z G

2 b

(13)

With the local microscopic heat-flux rate represented by Q/As , the

heat- transfer area from the specimen is represented by dt (d - average grain

size; t - width of the dislocation slip band), and - ?rd 2 /4 and - t

(-10 ^m) . Using these relations with Eq. 13, we obtain

As 4

a

2.a . Ud

d l

G

2 b

(14)

The average dislocation velocity T3 was estimated by Alden21 to be about

10” 6 m/s after the yield strength is reached at 295 K. Suzuki 22 related

dislocation velocity to exp B (a — a 0 ) with B a constant and o 0 related to

the athermal stress contribution. More commonly, Vreeland23 and many others

considered that E — am

, with the exponent, m, having values up to 40 for

face-centered cubic metals. Alden used E - ac/r 0 with ac an effective

stress containing a network-hardening term subtracted from the applied

stress and r 0 a viscous drag constant. From Eq. 14 and u - 10~ 6 m/s, G = 80

GPa at 4 K, and b - 3 x 10“ 1£> m; Q/As equals 4.6 x 10 3 W/m2. This heat-flux

rate is very close to the critical value for triggering the transition from

nucleate -boiling to film-boiling heat transfer (Fig. 5).

The local internal specimen heating from moving dislocations may also

be estimated using the model of Eshelby and Pratt. 11 To use Eq. 2 to

calculate the local temperature increase, we consider that a - D - 10 -6

m/s, b - 3 x 10” 10 m, and that at 4 K, k - 1.7 x 10” s m

2 /s and A - 0.5

102
W/(m-K). The number of moving dislocations, n, is again related to the

strain-hardening component of the flow stress as n - Pm' Ai and Pm “ cr£/(Gb) 2 following arguments previously presented. Here, represents the area

encompassing the moving dislocations and is taken as jrd 2 /4 (d - 10“ 4 ), the

equivalent of a slip band extending through a large grain. With these

assumptions and data, AT equals 47 K, a high value. Ideally, as the

temperature rises, A increases, and k remains approximately constant,

resulting in a reduced AT.

These models and calculations suggest that (1) at the microscopic

level, significant local temperature increases are achievable from moving

dislocations in low- conductivity materials (at low temperatures) and (2) at

low temperatures in helium gas or liquid, solid or surface heat- transfer

rates are not sufficient to keep up with the heat generated by the moving

dislocations

SUMMARY

At the strain rates at which we observed initiation of discontinuous

yielding, the calculated macroscopic specimen heating was not large enough

to trigger the transition from nucleate to film boiling and thus induce a

sudden, large temperature increase in the specimen. However, on a microscopic scale, we have shown that significant local temperature rises

(~ 47 K) may be induced by motion of multiple dislocations at 4 K in

austenitic steels. The heat flux from moving dislocations in surface grains

is of the same magnitude as the critical heat- transfer flux for the

transition from nucleate to film boiling. This suggests that the transition

in initiation stress ( a and strain (e^) for discontinuous yielding is

caused by local changes in the liquid heat- transfer mechanism. It also

suggests that the lower a

and values associated with tests in gaseous

helium at 4 K are directly attributable to the difference in heat- transfer

coefficients in helium liquid (nucleate boiling) and helium gas.

These conclusions imply that, after sufficient flow stress is attained,

the deforming specimen contains a multitude of local hot spots (slip bands).

As the specimen continues to deform, Q/A eventually increases beyond

(Q/A) max and triggers the transition to film heat transfer, abruptly changing the local temperature in the disturbed region.

The constancy of dislocation- interaction stress, for the initiation

of discontinuous yields for a large number of austenitic steels suggests

that a minimum number of dislocations are required to initiate the local

rise in temperature. This minimal dislocation density is independent of

stacking- fault energy, type and number of solutes, and yield- strength level

since a wide variety of austenitic steels are represented in the data. This

condition implies a local heating process that depends on dislocation

generation, rather than on pile-up stresses that are more dependent on the

above material parameters

103
ACKNOWLEDGMENTS

This work was partially supported by the Office of Fusion Energy,

U.S. Department of Energy. We thank V. D. Arp for essential consultations.

L. A. Delgado and R. P. Walsh carefully conducted most of the tensile

measurements

.

REFERENCES

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helium, in: "Advances in Cryogenic Engineering - Materials,"

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2. D. T. Read and R. P. Reed, Heating effects during tensile tests of AISI

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Toughness Testing of Low Temperature Structural Materials," National

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.

4. W. G. Dobson and D. L. Johnson, Effect of strain rate on measured

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low temperature instability of plastic flow, Acta Metall. 30:385-394

(1982).

10. K. Shibata, H. Sakamoto, K. Fujita, and T. Fujita, Computer simulation

of serration near liquid helium temperatures in: "Advances in Cryogenic Engineering — Materials," vol. 34, Plenum, New York (1988),

pp. 217-223.

11. J. D. Eshelby and P. L. Pratt, Note on the heating effect of moving

dislocations, Acta Metall. 4:560—562 (1956).

12. R. W. Armstrong, C. S. Coffey, and W. L. Elban, Adiabatic heating at a

dislocation pile-up avalanche, Acta Metall. 30:2111—2116 (1982).

13. G. A. Malygin, Local heating in dislocation ensembles at low temperatures, Sov. J. Low Temp. Phys. 5(11) : 633—637 (1979).

14. R. P. Reed and R. P. Walsh, Tensile strain-rate effects in liquid

helium, in: "Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures — X," NBSIR 87-3067, National Bureau of

Standards, Boulder, Colorado (1987), pp. 9—30.

15. R. P. Reed and R. P. Walsh, Low temperature tensile characteristics of

copper and aluminum, in: "Materials Studies for Magnetic Fusion

104
16 .

Energy Applications at Low Temperatures — IX," NBSIR 87-3050,

National Bureau of Standards, Boulder, Colorado (1986), pp. 107—140.

E. G. Brentari, P. S. Giarratano, and R. V. Smith, "Boiling Heat Transfer for Oxygen, Nitrogen, Hydrogen, and Helium," NBS Technical Note

317, National Bureau of Standards, Boulder, Colorado (1965).

17. W. B. Bald and T-Y. Wang, The nucleate pool boiling dilemma, Cryogenics

16:314-315 (1976).

18. R. M. Fand and J. Brucker, A correlation for heat transfer by natural

convection from horizontal cylinders that accounts for viscous

dissipation, Int. J. Heat Mass Trans. 26:709—726 (1983).

19. T. H. K. Frederking, Cryogenic heat transfer: He 4 Kapitza conductances

including phase change effects, in: "Heat Transfer 1982," vol. 1,

Hemisphere Publishing Corp., Washington (1983), pp 323-342.

20. L.-H. Lin and T. H. K. Frederking, Influence of the Kapitza resistance

on nucleate boiling of liquid helium, Lett. Heat Mass Trans. 9:473-

478 (1982).

21. T. H. Alden, Theory of mobile dislocation density: application to the

deformation of 304 stainless steel, Metall. Trans. A 18A: 51-62

(1987).

22. T. Suzuki, Dislocation motion and yield stress in pure copper and its

dilute alloys, in: "Dislocation Dynamics," McGraw-Hill, New York

(1968), pp. 551-572.

23. T. Vreeland, Dislocation velocity in copper and zinc, in: "Dislocation

Dynamics," McGraw-Hill, New York (1968), pp. 529—549.

105
-
INFLUENCE OF INTERSTITIAL CONTENT

ON FRACTURE TOUGHNESS

P.T. Purtscher and R.P. Reed

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado

The effects of interstitial content on tensile properties,

martensitic transformation, and inclusion content are examined in

relationship to the measured fracture toughness of austenitic

stainless steels at cryogenic temperatures. Metallographic

sections show that void nucleation occurs ahead of a crack only

after the applied J exceeds a large fraction of the measured

toughness, Jjc

. A stress intensity factor is calculated for void

nucleation from a simple model that considers the yield strength

and inclusion spacing.

INTRODUCTION

Increasing the interstitial carbon and nitrogen content increases the

yield strength of austenitic steels at low temperatures. The high strength

combined with high elastic modulus and good thermal properties make these

steels attractive for use in cryogenic structures. However, to build an

efficient, yet safe structure, high fracture toughness is required.

Experience 1 has shown that there is an empirical relationship between

yield strength (ay) and fracture toughness (Kj c ) for austenitic steels at

4 K:

KIc - 500 - 0.3a

y ±20 MPa^m [1]

A similar relationship is found for martensitic steels, aluminum, and

titanium alloys at ambient temperatures. 2

In practice, we know that other factors such as the inclusion content 3 * 4 and strain- induced martensitic transformation (TRIP effect) 6

" 7 influence

ductile fracture. Research has focused on ductile fracture characteristics

void nucleation, growth, and coalescence in the complex stress state ahead

of a crack. Most fundamental work has followed the Rice and Johnson approach8

" 10 of relating the toughness to growth of voids ahead of the crack,

assuming that void nucleation is relatively easy and contributed little to

* Contributed to the proceedings of international conference on High

Nitrogen Steels.

107
the magnitude of fracture toughness. Recently, we presented a model for estimating the stress intensity factor required to nucleate microvoids ahead

of a crack and argued that void nucleation is significant in ductile

fracture of austenitic steels at low temperatures. 11

In this paper, we present new mechanical test data at 76 and 4 K for a

series of seven austenitic stainless steels with varying carbon and nitrogen

contents. We examine the crack path, microstructure, and void growth at

loads near Jjc , using the light microscope and scanning electron microscopy

(SEM) . The purpose of this work is to show how the interstitial content of

austenitic steels influences fracture toughness.

MATERIALS AND PROCEDURES

The base composition of the seven steels is Fe-18Cr-10Ni-l.5Mn-0.022P0 . 017S-0 . 54S1-0 . 2Mo-0 . 2Cu. The interstitial content for each alloy is different and is shown in Table 1. All specimens were annealed at 1423 K for

2 h and water quenched.

The inclusion content of each steel was determined by examining a mechanically polished section transverse to the rolling direction in the SEM.

Inclusions were counted at two different magnifications by a computer-aided

technique that measured the average diameter, area, and shape factor for

each inclusion.

The tensile tests of all the steels were performed according to a newly

proposed standard for testing in liquid helium and are described in detail

by Reed and Simon. 12 The strain- induced transformation was monitored during

tensile testing at 76 K of steels 2, 5, and 8. The martensite content was

measured magnetically. Again, the details are described by Reed and Simon. 12

Fracture toughness tests were performed on all steels at 76 K and on

steels 2, 3, and 8 at 4 K. Compact specimens, 25 mm thick, were machined

from the annealed plates in the T-L orientation. The test procedures

followed the ASTM E 813-87 standard. Additional tests on #2 and 8 at 76 K

were interrupted at an applied J level < JjcTable 1 . Interstitial Content

Table 1. Interstitial Content

Alloy C(wt%) N(wt%) C+N(at%)

2 0.058 0.039 0.45

3 0.089 0.039 0.69

5 0.067 0.120 0.88

6 0.094 0.120 0.95

7 0.030 0.240 1.29

8 0.057 0.240 1.38

9 0.085 0.240 1.47

Base composition is Fe-18Cr-lONi.

108
*

CO

z

o

CO

Z>

—I

o

o

Figure 1. Histogram of the number and volume of inclusions for different

average diameter inclusions.

Following the toughness testing, all specimens were sectioned in half;

one half was fatigue post-cracked to reveal the fracture surface and the

other half was polished so that the cross section from the center of the

specimen could be examined in the SEM.

RESULTS

The majority of inclusions have an average diameter of < 1 pm and are

roughly spherical. The average spacing between inclusions at 400X was 46

to 53 pm and a total area fraction was about 0.0015. The spacing at 2000X

dropped to about 25 pm because numerous small inclusions were detected.

Figure 1 shows a histogram of the number and volume associated with

inclusions of different average diameters. The volume was estimated by

assuming an average diameter for each bracket, calculating the volume of one

inclusion, and multiplying by the number of inclusions in the given bracket.

A complete summary of the tensile results are found in a separate

paper. 12 The fracture toughness results are plotted as a function of

tensile yield strength in Figure 2. The results are similar to the previous

trend. 1

Figure 3 shows normalized values of the percent martensite content

plotted as a function of interstitial content for various amounts of applied

strain in the tensile test. For steel 2, most of the transformation occurs

between 10 and 30% elongation. The maximum martensite content was about

85%. For alloy 5, a similar trend is observed, but the maximum value is

reduced to 60%. In alloy 8, the martensite content increases with strain,

but the amount of transformation is smaller and more uniform. The trends at

4 K are similar to 76 K, but that the maximum amount of martensite which

forms at 4 K is slightly lower. 13

109
YIELD STRENGTH, o-v (MPa)

Figure 2. Fracture toughness vs. yield strength data at 76 and 4 K compared

to previous 4 K data

.

1

Figure 4 shows the microstructure ahead of the crack in alloy 2 tested

at 76 K to an applied J - 0.35 Jjc

. The heavily deformed areas are at

angles to the plane of maximum stress that lies directly ahead of the crack.

The dark etching areas contain strain- induced martensite. Little martensite

is present directly ahead of the crack; this implies that the plastic strain

in this region was relatively low. Finite element calculations from

different sources 14 ’ 15 indicate that the stresses directly ahead of the

crack should have a maximum value of 3 to 4 times the tensile yield stress.

For alloy 8 tested at 76 K (applied J # - 0.8 Jjc)» the load- displacement

curve was nearly linear-elastic and there was little plasticity evident at

the crack tip.

110
0 0.5 1.0 1.5

(C+N), at.%

Figure 3. Martensite transformation vs. interstitial content as a function

of strain at 76 K for alloys 2, 5, and 8.

The fracture surfaces from the toughness specimens tested at 76 K consisted mainly of equiaxed dimples, about 50 pm in diameter, that nucleated

from the larger inclusions in the matrix. Smaller dimples arranged in

planar arrays to connect the large dimples in three dimensions in a manner

that may be described as void sheet coalescence. Higher percentages of the

small dimples were found on specimens from alloys 2 and 3. Figure 5 shows a

representative fractograph from alloy 8 at 76 K. At 4 K, the fracture surfaces were similar, but there is a larger percentage of smaller dimples

between the 50 pm diameter dimples. No stretched zone was observed between

the fatigue precracked surface and the dimpled rupture features that could

be related to the measured toughness. 16

Ill
Figure 4 Etched microstructure of the area around the blunted crack tip in

alloy 2. The test was interrupted at an applied energy = 0.35 Jjc

.

112
I I 20 yin

Figure 5. Fracture surface from toughness specimen of alloy 8 tested at

4 K.

Examination of sectioned specimens that had been loaded to applied

energies less than Jj c showed that no voids could be found ahead of the

crack. For alloy 2, the test was stopped at an applied J equal to

approximately 35% of Jj c ; for alloy 8, the applied J was 80% of the critical

value. The direction of crack growth in the specimens tested in the normal

fashion was confined to the region of high stress, low strain ahead of the

crack. Figure 6 shows a sequence of blunting and crack growth that illustrates these observations in alloy 8 tested at 76 K.

DISCUSSION

The interstitial content influenced the tensile yield strength and the

strain- induced martensitic transformation as expected, 17 but the inclusion

content was found to be similar in all seven steels. These results are

typical of all the reported cryogenic mechanical properties of austenitic

steels

1

8

Extensive strain- induced transformation occurs during testing depending

upon the test temperature and interstitial content. However, the geometries

of the uniaxial tension and fracture mechanics specimens produce different

stress and strain patterns. Transformation does not occur prior to void

nucleation at the inclusions that will eventually form the fracture surface

in the toughness test. Apparently, the local stress and strain that produce

void nucleation are not influenced directly by strain- induced martensite

formation.

113
Figure 6 SEM micrographs of the polished cross section from crack-tip

regions in alloy 8, a) test stopped at applied J - 0.8 J Ic and b) test

stopped at applied J > Jjc

.

114
Observations of the sectioned toughness specimens suggest that the

applied energy associated with nucleation of voids accounts for a large

fraction of the measured fracture toughness . The growth of voids from these

larger inclusions and coalescence by nucleation and growth of smaller voids

apparently requires little additional applied energy. Therefore, growthcontrolled models for ductile fracture will not be very precise.

LeRoy et al. 18 have described void nucleation in terms of the critical

interfacial strength between the matrix and inclusions (s c ) and assuming

s c " SH U + ° [2]

where a — the applied (macroscopic) stress, c - plastic strain and sjj — a

material constant so that the product {s^Vc) portrays the microscopic stress

from dislocation pileups. A critical strain can be calculated in terms of

the constants, s c and sjj, and the applied stress. The critical strain can

then be used for the fracture strain (cf) in an expression like

eq. 3 from Ashby20 to form eq. 4

KIc - (2xOyT*Ee f)h [3]

Knuc -

(

S

|r

sH

5gy

) (

4 5oyEI.j

H

[4]

where E - Young's modulus, r* - a characteristic distance like the inclusion

spacing, Knuc - stress intensity factor for nucleation, L - r*, and

representing the stress in front of the crack tip, a - 3.5<7y.

Values for sjj and s c were estimated in previous work1

to be about 2400

and 5000 MPa. Using a value of E - 200 GPa and L - 50 nm, we can calculate

a value of for each test temperature and composition. The calculated

values are compared to the measured values and plotted in Figure 7 as a

function of interstitial content. The best agreement occurs for the 4 K

data and higher interstitial contents, indicating that nucleation accounts

for a larger share of the measured fracture toughness at lower temperature

and higher strengths.

CONCLUSIONS

1. Increasing the interstitial content raises the yield strength, makes

void nucleation ahead of the crack easier, and reduces fracture toughness of austenitic stainless steels at low temperatures.

2. Increasing the interstitial content reduces the extent of the straininduced transformation. The transformation does not appear to affect

the local stress and strain at the inclusions which form the fracture

surface

3. Increasing the interstitial content does not affect the inclusion size

and distribution.

115
Figure 7. Toughness ratio (predicted nucleation stress intensity factor

divided by the measured toughness) plotted vs. the interstitial

content

.

ACKNOWLEDGMENTS

This work was partially sponsored by DOE, Office of Fusion Energy. The

authors thank R.P. Walsh, D.A. Shepherd, and L. Delgado for their help in

obtaining the experimental data.

REFERENCES

1. R.L. Tobler, D.T. Read, and R.P. Reed: in Fracture Mechanics

Thirteenth Conference . ASTM STP 743, R. Roberts, ed. , American Society

for Testing and Materials, Philadelphia, 1981, p. 350.

2. R.J.H. Wanhill: Eng. Fract. Mech. , vol. 10, 1978, p. 337.

3. J.C. Lautridou and A. Pineau: Eng. Fract. Mech., vol. 15, 1981,

p. 1259.

4. A.D. Wilson: J. Eng. Matl. Tech., vol. 101, 1979, p. 265.

116
5. E.R. Parker and V.F. Zackay: Eng. Fract. Mech. , vol. 5, 1973, p. 147.

6. D. Fahr: Met. Trans., vol. 2A, 1971, p. 1883.

7. S. D. Antolovich and B. Singh: Met. Trans., vol. 2A, 1971, p. 2135.

8. J.R. Rice and M. A. Johnson: in Inelastic Behavior of Solids . M.F.

Kannimen, et al., eds., McGraw-Hill, New York, 1970, p. 641.

9. R.O. Ritchie and A.W. Thompson: Met. Trans., vol. 16A, 1985, p. 233.

10. W.M. Garrison, Jr.: Met. Trans., vol. 17A, 1986, p. 669.

11. P.T. Purtscher, D.T. Read, and R.P. Reed: Fracture Mechanics

Perspectives and Directions (Twentieth Symposium) . ASTM STP 1020, R.P. Wei

and R.P. Gangloff, Eds., American Society for Testing and Materials,

Phialdelphia, 1989, pp. 443-446.

12. R.P. Reed and N.J. Simon: in these proceedings.

13. R.P. Reed and C.J. Guntner: Trans. AIME, vol. 230, 1964, p. 1713.

14. R.M. McMeeking: J. Mech. Phys. Solids, vol. 25, 1977, p. 357.

15. C.F. Shih: J. Mech. Phys. Solids, vol. 29, 1981, p. 305.

16. P.T. Purtscher: J. Test. Eval., vol. 15, 1987, p. 296.

17. R.P. Reed: in Materials at Low Temperature . R.P. Reed and A.F. Clark,

eds., American Society for Metals, Metals Park, Ohio, 1983, p. 295.

18. J.W. Morris, Jr. and E.N.C. Dalder: J. Metals, vol. 37, 1985, p. 24.

19. G. LeRoy, J.D. Embury, G. Edwards, and M.F. Ashby: Acta Met., vol. 29,

1981, p. 1509.

20. M.F. Ashby: Prog. Mat. Sci.;Bruce Chalmers Anniversary Volume, p. 1.

117

METALLOGRAPHIC STUDY OF THE CRACK- TIP REGION FROM FRACTURE

MECHANICS SPECIMENS OF AUSTENITIC AND FERRITIC STEELS'4

"

P.T. Purtscher*, R.P. Reed*, and D.K. Matlock**

Fracture and Deformation Division

*National Institute of Standards and Technology Division

Boulder, Colorado

**Advanced Steel Processing and Products Research Center, Colorado

School of Mines

Golden, Colorado

Metallographic sections of blunted and growing cracks in

steels with different strain-hardening capacities are observed in

the light and scanning electron microscopes. The strain distribution around the crack- tip region is inferred from the etching

behavior and microhardness readings. The direction of initial

growth is approximately straight ahead of the fatigue precrack in

the high strain-hardening austenitic steels and at an angle of

about 45° in the- low strain-hardening ferritic steel. These

observations are discussed in terms of different criteria for

ductile fracture in fracture mechanics specimens.

INTRODUCTION

The absence of stretch zones on the fracture surface of austenitic

steels tested at cryogenic temperatures 1 indicates that the strain distribution at the crack tip must be significantly different from the strain around

cracks in ferritic steels. The stress and strain distributions around a

crack in fracture mechanics specimens affect the fracture criteria used in

modeling fracture toughness. McMeeking2 suggested, based on his finite

element calculations, that if the crack grows straight ahead of the tip, a

critical-stress fracture criterion is appropriate (example in ref. 3); on

the other hand, angled crack- growth indicates a critical plastic strain

criterion (examples in ref. 4 and 5).

Many criteria have been proposed to describe ductile fracture in

fracture mechanics specimens. One approach6 relates fracture toughness to

the growth of voids ahead of the crack tip, assuming that void nucleation

contributes little to the measured toughness. Others 7 * 8 have followed this

approach, but incorporated microstructural features to describe the local

strain and the characteristic distance for fracture. In contrast, an

expression for the stress intensity factor for void nucleation has been

developed and shown to compare favorably to fracture toughness in highstrength austenitic steels. 8

1 Contributed to proceedings of MRS Advanced Materials meeting in Tokyo

,

June, 1988.

119
In this paper, the crack- tip regions in fracture mechanics specimens of

austenitic and ferritic steels are examined and the initial crack growth

directions are determined. The observations are discussed with respect to

the strain-hardening capacities of the steels, finite element results, and

different fracture criteria.

MATERIALS AND PROCEDURES

Two austenitic steels tested at 76 K and a ferritic steel tested at

233 K, are used in this study. Alloy #1 is a low strength, high strainhardening austenitic stainless steel (similar to AISI 304); alloy #2 is a

higher strength, austenitic stainless steel (high nitrogen version of AISI

304) ; alloy #3 is a ferritic steel (similar to ASTM A533B) with a yield

strength in between that for alloys #1 and 2. The chemical compositions and

fracture toughnesses of the steels are listed in Table 1 and the true

stress—true strain curves are shown in Figure 1. A complete analysis of the

mechanical properties of these steels and test procedures can be found

elsewhere

.

1

0

"

1

1

Table I. Steel compositions (wt. %) and fracture toughness.

Alloy Cr Ni Mn P S Si Mo C N KIc (MPaym) T(K)

1 18.3 10.1 1.6 0.021 0.019 0.5 0.2 0.06 0.04 375 76

2 20.3 10.0 1.4 0.024 0.016 0.5 0.2 0.06 0.24 270 76

3 0.1 0.8 1.3 0.007 0.008 0.2 0.5 0.19 — 256 233

Standard and interrupted fracture toughness tests (i.e. applied J, J a , less than the fracture toughness, Jjc ) were performed on side-grooved,

compact specimens of alloys #1 and 2 in liquid nitrogen (76 K) . A similar

specimen from alloy #3 was tested in cooled alcohol (233 K) where a ductile

crack initiated at the fatigue crack tip, but the specimen ultimately failed

by cleavage. The austenitic specimens were loaded to a crack extension of

about 1.5 mm, then unloaded, warmed to room temperature and sectioned along

the centerline perpendicular to the crack plane. One half was fatigue postcracked to reveal the extent of crack growth during the test while the other

half was polished to observe the crack- tip geometry. The fracture surface

from the broken specimen of alloy #3 was observed in the scanning electron

microscope before the two halves were glued together with epoxy for sectioning. Diamond pyramid hardness readings (DPH) and etching solutions 12

were used to delineate the strain-hardened regions around the crack tips.

120
4.0

CO

CL

0

3.0

' ' v #1, T = 76 K

/ /

#2, T = 76 K

-x

#3, T= 233 K

0.4 0.8 1.2

TRUE STRAIN

1.6

Figure 1. True -stress -strain curves for the three alloys. Dotted lines

represent post necking extrapolation to measured (indicated by X)

stress -strain point at fracture.

OBSERVATIONS

No stretch zones formed on the fracture surface of the compact specimens from alloys #1 and 2. In contrast, a small, well-defined stretch zone

is found on the fracture surface of alloy #3 between the fatigue precrack

features and the dimpled rupture features of the overload region. These

observations are similar to that shown in ref. 1.

121
Figure 2. Hardness gradients from grip to necked region in tensile tests

and around the crack tip for the three alloys

.

Figure 2 shows the gradients of microhardness readings from tensile and

fracture toughness test specimens. The range of DPH measured in the tensile

test defines the strain-hardening capacity (SHC) of each steel. Alloys #1

and 2 have a high SHC while alloy #3 has a lower SHC. A large gradient in

the DPH values was observed around the crack tip for alloy #1 with a DPH

close to the crack tip that approaches the maximum DPH measured in the

necked region of the tensile specimen. For alloy #3 which had the lowest

SHC in the tensile test, the DPH values around the crack tip are more

uniform and spread over a larger area. The maximum DPH near the tip is

similar to the maximum measured in the tensile test.

122
Figure 3 shows representative dimpled rupture features from both the

austenitic and ferritic steel samples. Alloys #1 and 2 (Fig. 3a) exhibited

a range of dimple morphologies where the dimples formed primarily at MnS

inclusions. Alloy #3 (Fig. 3b) displayed smaller and more uniform dimples

which nucleated at the carbides that form on tempering the steel.

Figure 3. Fracture surfaces from compact specimens (a) alloy #2 tested at

76 K and (b) alloy #3 tested at 233 K.

123
The sectioned crack from alloy #1 interrupted at J a - 0.35 Jjc is shown

in Fig. 4. The stain etch (Fig. 4a) shows that there is a sharp gradient in

the strain- induced martensite concentrated at 45° to the crack plane. In

addition, selected microhardness readings indicate that the strain hardening

is concentrated in the dark-etching bands. With a different etch and higher

magnification (Fig. 4b), a small crack can be seen directly ahead of the

blunted tip.

Figure 4. Section of crack tip from alloy #1 at J a - 0.35 Jj c (a) stain

etch 12 and (b) oxalyic acid etch.

124
A specimen of alloy #2 stopped at J a - 0.8 Jj c , Fig. 5a, exhibited a

smaller crack opening displacement (COD or 8 ) (in comparison to alloy #1).

DPH readings and the stain etch indicated that the region of strain hardening were proportionately smaller, but retained a shape similar to that

shown in Fig. 4. No crack was present at the tip, but the blunted shape

extended directly ahead and does not show shear crack formation along the

strain hardened regions at 45° to the crack plane.

A section through the growing crack in alloy #2, J a > Jj c , i- s shown in

Fig. 5b. Few voids have opened around the growing crack and the DPH profile

(Fig. 2) indicates that the strain hardening is concentrated in a zone that

is smaller than the plastic zone.

In Fig. 6 from alloy #3 (J a > Ji c ), the initial growth has a zig-zag

appearance similar to crack growth associated with slip- line fields. 6 The

geometry of the plastic zone in A533B steel has been shown 13 to match theoretical predictions 2 with dimensions in the plane of the crack equal to the

COD, and equal to twice the COD at 45°. Again, void nucleation and growth

are confined to a zone smaller than the plastic zone determined by the DPH

profile (Fig. 2)

DISCUSSION

The straight-ahead growth in the austenitic steels is interpreted to

be controlled by the peak stress (3 to 4 times the tensile yield stress

according to the finite element predictions 2

) ahead of the crack while the

angled crack growth in the ferritic steel is interpreted to be controlled by

strain. In both cases, few voids are observed to have formed away from the

growing crack in the plastic zone. The crack-growth behavior and the appearance of the crack- tip region (stretch zone formation) in the two types

of steel can be explained by the strain hardening.

When the crack opens in alloys #1 and 2 (higher SHC) , the crack- tip

strength increases from localized strain hardening, and a relatively small

plastic zone develops directly ahead of the blunted tip (shown by the micro

hardness gradients and the etching behavior) . The matrix which will

eventually fail and form the fracture surface does not flow significantly

until voids nucleate at the larger MnS inclusions. The increased flow

stress due to constraint and the low test temperature produce void nucleation on the crack plane where stress is highest, directly ahead of the tip.

After nucleation, the matrix can flow around the particles leading to void

growth and coalescence. Again because of the low test temperature, the flow

stresses are high and void growth is minimal. No stretch zone is observed

because the material on the fracture surface was not strained appreciably

prior to void nucleation.

In contrast, when the crack opens in alloy #3 with the lower SHC, the

crack tip has a lower strength and the strain hardening is distributed over

a larger plastic zone (evidenced by the microhardness data) . The plastic

zone extends ahead of the crack, but the highest strains are concentrated in

bands at 45° to the tip where voids nucleate at small tempered carbides.

125
Figure 5. Crack- tip sections from specimens of alloy #2: (a) J a - 0.8 Jj c shows no void nucleation ahead of the blunted tip and (b) J a >

Jjc showing the straight-ahead growth.

126
Figure 6. Section of crack tip in alloy #3 with J a > Jj c showing the angled

path of initial crack growth.

The extent of growth is limited by the spacing of the particles; that is,

little additional COD is required for the voids to grow and coalesce to

produce the angled crack growth. The stretch zone forms because the

material on the fracture surface was strained appreciably prior to void

nucleation.

In both types of steels, void nucleation was limited to a process zone

which can be defined by the observations to have a diameter approximately

equal to the plastic COD. Conditions for void nucleation are only met in

this zone where the stress and strain gradients are large. The measured

fracture toughnesses (Table 1) and a general failure criterion must be

related to the effect of stress and strain on both void nucleation and

growth within the process zone. The contribution of void growth is small

here because of the short spacing between carbides in alloy #3 and the low

test temperature for alloys #1 and 2. The size of the process zone should

be related to the COD, which determines the scale over which the enhanced

forces act, and the size and spacing of the microstructural features (second

phase particles), which determine the resistance of the steel to fracture.

SUMMARY and CONCLUSIONS

All of the steels tested for this investigation fail in a ductile

manner, and the measured fracture toughness values are relatively high. A

stretch zone is observed on the fracture surface of a ferritic steel tested

at 233 K where the initial crack growth is at 45° to the blunted tip. No

127
stretch zone is observed on the fracture surface of the austenitic steels

tested at 76 K where the initial crack growth is straight -ahead. These

observations have led to the following conclusions:

1 . Stretch zones form on the fracture surface when the material is

strained significantly prior to void nucleation.

2. In the austenitic steels (high SHC) where no stretch zone forms

and the crack grows directly ahead of the blunted tip, fracture

toughness is controlled by the peak stress ahead of the crack

tip.

3. In the ferritic steel (low SHC) where a stretch zone forms on the

fracture surface and an angled crack- growth path is observed,

fracture toughness is controlled by the maximum strain which

occurs at 45° to the crack tip.

4. In both cases, void nucleation was observed in only a limited

region around the crack. The conditions for nucleation, and

similarly for fracture toughness, are best described by a

function of both stress and strain over a distance that is

related to the COD and the size and spacing of void- nucleating

particles

.

ACKNOWLEDGEMENTS

The DOE Office of Fusion Energy provided partial support of this work.

The authors wish to thank Jurgen Heerens from the GKSS -Research Center,

Geesthacht, West Germany for providing machined specimens of the ferritic

steel

.

REFERENCES

1. P.T. Purtscher, J. Test. Eval., vol. 15, 1987, p.296.

2. R.M. McMeeking, J. Mech. Phys . Solids, vol. 25, 1977, p. 357.

3. S.P. Rawal and J. Gurland, in Proceedings of Second International

Conference on Mechanical Behavior of Materials . Federation of Materials

Societies, Dearborn, MI, 1976, p. 1154.

4. C.D. Beachem and G.R. Yoder, Metall. Trans., vol. 4A, 1973, p. 1145.

5. F. Zia-Ebrahimi , D.K. Matlock, and G. Krauss, Scripta Metall., vol. 16,

1982, p. 987.

6. J.R. Rice and M.A. Johnson, in Inelastic Behavoir of Solids , edited by

M.F. Kanninen, W.F. Adler, A.R. Rosenfield, and R.I. Jaffee, McGrawHill, New York, 1968, p. 641.

7. R.O. Ritchie and A.W. Thompson, Metall. Trans., vol. 16A, 1985, p. 233.

128
8. W.M. Garrison, Jr., Metall. Trans., vol. 17A, 1986, p. 669.

9. P.T. Purtscher, D.T. Read, and R.P. Reed, Fracture Mechanics: Prosoectives and Directions (Twentieth Symposium') . edited by R. Wei and R.P.

Gangloff, American Society for Testing and Materials, Philadelphia,

1989, pp. 433-446.

10. P.T. Purtscher and R.P. Reed, to be published in Proceedings of High

Nitrogen Steels '88 .

11. J. Heerens, K. Schwab le, and A. Cornec , in Fracture Mechanics:

Eighteenth Symposium . ASTM STP 945, edited by D.T. Read, and R.P. Reed

American Society for Testing and Materials, Philadelphia, 1988, p.374.

12. R.P. Reed, Acta Metall., vol. 10, 1962, p. 865.

13. R.H. Van Stone, T.B. Cox, J.R. Low, Jr., and P.A. Psioda, Int. Met.

Rev., vol. 30, 1985, p. 157.

129
*

^ . inS “ •-'& >»-<>*<» i
TENSILE, FRACTURE, AND FATIGUE PROPERTIES OF NOTCHED ALUMINUM

ALLOY SHEETS AT LIQUID NITROGEN TEMPERATURE

R.L. Tobler, R.P. Walsh and R.P. Reed

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado

J.K. Han

Korea Advanced Institute of Science and Technology

Seoul, Korea

L. Ma

Institute of Metal Research Acedemia Sinica

Shenyang, People's Republic of China

Notched sheet specimens of four high strength aluminum alloys

were tested to evaluate their potential for a Superconducting

Super Collider application. Alloys 7075-T6, 7475-T761, 2219-T87,

and 2090-T8E41 were tested in liquid nitrogen at 76 K, in tensiontension fatigue, at a stress ratio of 0.1. The alloys are

compared in terms of their static tensile and fracture properties,

and fatigue resistance for a life of 4 x IQ 4 cycles.

INTRODUCTION

Notched key collars are to be used in the Superconducting Super

Collider (SSC) dipole magnets (1). The key collars are locked in position

on cryogenic magnetic coils so the magnetic coils can be prestressed when

the keys are pressed in. Safe structural design is critical because the key

collars are highly stressed repeatedly at 4 K. A fatigue life of 4 x 10 4 cycles is anticipated. So far, alloy 7075-T6 has been selected as the

construction material. Preliminary analysis and experiments support the

adequacy of 7075-T6: model coils were energized 3500 times to the full

operating field strength of 6.6 T; they were also energized several times

to 9 T, at 1.8 K, and no cracking was observed in the collars (1). In this

study, we report various mechanical property measurements enabling a preliminary comparison of the relative merits of 7075 -T6 and three competing

alloys at 76 K.

131
MATERIALS

Four commercial aluminum alloys were tested in their as -received

tempers: 7075-T6, 7475-T761, 2219-T87, and 2090-T8E41. Alloy 7475-T761 was

obtained in the form of a 2.10 mm thick clad sheet because unclad sheet was

not available from several suppliers. The other three alloys were obtained

in the desired form, 1.60 mm thick sheets, unclad. Nominal compositions in

weight percent are:

7075 — 5 . 6%Zn-2 . 5%Mg- 1 . 6%Cu-0 . 23%Cr , balance Al

7475 - 5 . 7%Zn-2 . 3%Mg- 1 . 5%Cu-0 . 22%Cr , balance Al

2219 - 6 . 3%Cu-0 . 3%Mn-0 . 18%Zr-0 . 10%V-0 . 06%Ti , balance Al

2090 - 2 . 86%Cu-2 . 05%Li-0 . 12%Zr , balance Al

Figure 1 shows the specimen geometries. Unnotched specimens for

routine tensile tests were fabricated with the loading axis transverse to

the principal rolling direction of the sheet. The notched specimens for

tensile and axial fatigue tests were also fabricated in the transverse

orientation. The square notch shown in Fig. 1 essentially duplicates the

notched condition of the actual SSC key collar: the radius, ligament size,

thickness, and notch depth of our specimens are nearly the same as for the

part used in service. The actual key collar notch has tapered sides with an

angle of 3° (1), but this taper has a negligable effect on the notch stress

concentration factor Kt , which is approximately 3.9 for both the key collar

and our specimens.

(Dimensions in mm)

102

-25

12.7

-*—2 5 —

4>

n 18.7

1

12.7

Notch radii: 0.5 1

Figure 1

.

Specimen geometries.

132
PROCEDURES

All specimens were ground and polished to remove surface markings from

the as-received sheets. The 7475-T761 material, originally 2.10 mm thick,

included superficial layers of alcladding 0.08 mm thick. These specimens

were ground equally on each side, not only to remove the cladding (typically, alloy 7072) but to reduce the thickness to more closely match the

thicknesses of the other three alloys. Specimens of the three other

(unclad) alloys were ground slightly on each side. Finally, all specimens

were electro-polished in a solution of 30% nitric acid and 70% methanol.

After polishing the specimens, we measured their dimensions and notch

root radii using a micrometer or a shadowgraph. The nominal dimensions are

given in Fig. 1. The notch radii were measured at both comers and on both

sides of each specimen, and the four readings were averaged. The minimum

values of the four readings for each specimen were 0.555 mm ± 7%. The mean

values of the four readings for each specimen were 0.600 mm ± 10%. The

specimen thicknesses varied with the extent of grinding and polishing;

typical values were: 7075-T6, 1.53 mm; 7475-T761, 1.81 mm; 2219-T87,

1.61 mm; 2090-T8E41, 1.45 mm.

Conventional tensile tests were performed at a crosshead rate of

0.5 mm/min and a corresponding strain rate of 3.3 x 10" 4 s" 1

. The gage

length for elongation measurements was 29 mm. The unnotched tensile

specimen performed well for alloys 2090, 2219, and 7475, but during tests of

7075-T6 several specimens failed at the pin holes in the grips. To complete

the characterization of 7075-T6, we machined some modified specimens that

had a reduced gage-section width of 6.35 mm; this ensured failures in the

gage length.

Tensile tests of notched specimens were performed to measure the notch

tensile strength (NTS) at- 76 K. The first series of 2090-T8E41 specimens

used for this purpose had irregular notch tip radii, which caused low NTS

values and low fatigue resistance; the data for these specimens were

discarded and new specimens were machined.

The specimens were gripped and mounted in a cryostat for testing.

Liquid nitrogen surrounds the specimen for tests at 76 K. Fatigue tests

were conducted using servohydraulic machines in the load control mode with

10 or 25 kN load ranges. The minimum- to -maximum fatigue stress ratio R was

0.1. To begin a test, the mean load was applied at once, and the load

amplitude was adjusted while the frequency f was only 1 Hz. The frequency

was then gradually stepped up to 9 , 10, or 15 Hz in increments of 1 Hz while

the load amplitude was re -adjusted.

133
RESULTS

Tensile Properties

Table 1 lists the tensile test data for unnotched specimens of all

alloys tested in this study. With decreasing temperature between 295 and

76 K, the yield strengths YS and ultimate tensile strengths UTS increase

significantly while the values of tensile elongation EL and Young's modulus

E vary slightly. Such trends are typical for aluminum alloys and present no

surprises

.

Table 1. Tensile Test Results for Unnotched Specimens.

Alloy

Temp

.

K

YS

MPa

UTS

MPa

El

%

E

GPa

7075 295 504 589 16.5 61.2

-T6 295 499 589 17.3 65.1

avg - 502 avg - 589 avg - 16.9 avg -63.2

76 589 716 15.5 76.4

76 589 711 16.0 70.4

avg - 589 avg - 714 avg - 15.8 avg - 73.2

7475 295 462 514 17.3 65.9

-T761 295 458 515 16.8 66.6

avg - 460 avg - 515 avg - 17.1 avg -66.3

76 548 635 17.5 78.2

76 549 637 17.0 72.3

avg - 549 avg - 636 avg - 17.3 avg - 75.3

2219 295 398 475 12.8 69.9

-T87 295 396 475 12.4 65.7

avg - 397 avg - 475 avg - 12.6 avg -67.8

76 484 596 14.0 80.9

76 484 598 12.9 72.0

avg - 484 avg - 597 avg - 13.5 avg - 76.5

2090 295 488 528 12.0 77.2

-T8E41 295 487 527 12.1 70.8

avg - 488 avg - 528 avg - 12.1 avg - 74.0

76 552 637 10.5 - . 76 550 643 11.1 76.9

avg - 551 avg — 640 avg — 10.8

Note: The unnotched specimen diagrammed in Fig. 1 was used in

most tensile tests, but 7075 at 76 K required a modified specimen

(see text)

.

1 34
Fracture Resistance

Table 2 lists the tensile test data for notched specimens at 76 K.

Also shown are the notch- strength ratios (NSR - NTS/UTS) and notch-yield

ratios (NYR - NTS/YS). These ratios are widely used as qualitative

indicators of notch sensitivity and fracture toughness. In particular, it

has been shown that the NYR for round bar specimens generally correlates

with the plane-strain fracture toughness parameter, Kjc (2). The NYR's for

our alloys range from 0.76 to 0.90, and are always less than 1, indicating

that the alloys are notch sensitive. Figure 2, a plot of YS versus NYR for

the four alloys, demonstrates there is an inverse relationship such as would

be expected for parameters associated with strength and toughness.

As indicated by Fig. 2, the fracture resistance of alloy 2090-T8E41 is

not outstanding compared to the other alloys of this study at 76 K. This

came as a surprise since superior toughness at cryogenic temperatures has

been reported elsewhere for alloy 2090 (4—6). This may imply that rankings

based on qualitative NYR data for thin specimens (our study) yield different

conclusions from quantitative rankings based on Kjc values for thick

specimens (other studies) . Or perhaps the disparate findings reflect the

fact that the fracture toughness of alloy 2090-T8E41 at cryogenic temperatures depends on metallurgical factors (such as extent of recrystallization,

degree of texturing, and distribution of precipitates) which can vary from

heat to heat.

Table 2. Notched Tensile Strengths (NTS), Notch- Strength Ratios (NSR),

Notch-Yield Ratios (NYR) , and Notched Fatigue Strength Ratios

(NSFR) at 76 K.

Alloy Specimen NTS

MPa

NSRa NYRb NFSRC

at 10 6 cycles

7475-T761 1 483

2 461

avg - 472 0.74 0.86 0.21

7075-T6 1 478

2 423

avg - 450 0.63 0.76 0.21

2219-T87 1 450

2 420

avg - 435 0.73 0.90 0.21

2090-T8E41 1 425

2 464

3 441

avg - 443 0.69 0.80 0.23

a

b

c

Obtained by dividing NTS by UTS

Obtained by dividing NTS by YS

Obtained by dividing NFS at 10 6 cycles by NTS.

135
As indicated by Fig. 2, the fracture resistance of alloy 2090-T8E41 is

not outstanding compared to the other alloys of this study at 76 K. This

came as a surprise since superior toughness at cryogenic temperatures has

been reported elsewhere for alloy 2090 (4—6) . This may imply that rankings

based on qualitative NYR data for thin specimens (our study) yield different

conclusions from quantitative rankings based on Kjc values for thick

specimens (other studies) . Or perhaps the disparate findings reflect the

fact that the fracture toughness of alloy 2090-T8E41 at cryogenic temperatures depends on metallurgical factors (such as extent of recrystallization,

degree of texturing, and distribution of precipitates) which can vary from

heat to heat.

Fatigue Life

Our fatigue data for notched specimens at 76 K are presented in Fig. 3.

Typically, the measurements consist of two tests at each of three stress

levels bracketting a life of 4 x 10 4 cycles. The average life at a given

stress is plotted in Fig. 3.

As a measure of fatigue resistance, we report the notched fatigue

strength, NFS; this is the maximum stress corresponding to fatigue failure

at a specified number of cycles. The NFS values were estimated from the

trend lines in Fig. 3. For 4 x 10 4 cycles at 76 K, the alloy ranking in

descending order is

2090-T8E41

7475-T761

7075-T6

2219-T87

NFS - 129 MPa,

NFS - 109 MPa,

NFS - 103 MPa, and

NFS - 102 MPa.

650

600 o 7075-T6

-

550 o

2090-T8E41

O 7475-T761

-

500

T = 76 K

1

O 221 9-T87

450

0.8 0.9 1.

NOTCH-YIELD RATIO, NYR

Figure 2. Yield strength versus notch-yield ratio for the four alloys tested

at 76 K.

136
CYCLES TO FAILURE, N

Figure 3. Notched specimen fatigue test results for the four alloys tested

at 76 K.

To compare the NFS values with NTS values for these alloys we define a

parameter called the notched fatigue strength ratio NFSR - NFS/NTS . We made

this comparison at 10 s cycles where the NFS estimates from Fig. 4 are most

reliable. As shown in Table 2, for a life of 10 s cycles, the NFSR's for all

four alloys are in close agreement at 0.21-0.23. We conclude that 2090-

T8E41 offers a higher fatigue resistance in the notched condition at 76 K,

but the differences among these high strength aluminum alloys can be

normalized to NTS values.

DISCUSSION

General Observations

Fatigue data at cryogenic temperatures are not abundant in the open

literature but some data that are available indicate that the face centered

cubic metals including aluminum alloys show some regular property trends.

McCammon and Rosenberg (3) observed that the conventional fatiguestrengths

of unnotched specimens at cryogenic temperatures tend to increase with

increasing ultimate tensile strength. Similarly, our data suggest that the

fatigue strength of notched specimens scales approximately with the notched

tensile strength, under certain conditions (that is, a life of about 10 6

cycles). Therefore the conventional tensile properties of smooth and

notched specimens bear some relationship to their fatigue resistances, and

the tensile properties are potentially useful for screening.

137
Alloy screening

Although the NTS is related to 1/YS (Fig. 2), there is not an exact

correlation, and the NFS apparently scales better with NTS than with 1/YS.

For four alloys, we found that the NFS is nearly a constant fraction (0.21-

0.23) of the NTS (Table 2); this implies that tensile tests of notched

specimens, which are relatively easy to perform, can help to predict the

fatigue resistance of notched specimens at 10 s cycles. More data are needed

to confirm the value of the NFS/NTS ratio as a screening parameter for

fatigue, and to ascertain the range of applicability for various alloys,

temperatures, and stress levels.

Alloy Selection

Many aluminum alloys that remain ductile and tough at low temperatures

can be used in cryogenic structures but different alloys have advantages and

disadvantages that influence material selection for specific applications.

In this study at 76 K, no one alloy is universally superior: 7075-T6 offers

the highest yield and tensile strengths in the unnotched condition, 7475-

T761 offers the highest strength in the notched condition, 2219-T87 offers

the highest notch-yield ratio (fracture toughness), and 2090-T8E41 offers

the highest notched fatigue strength for the cycle life of interest. The

material availability and cost must be considered as well as the mechanical

properties

.

Among the 7000 series alloys, 7075-T6 is readily available and widely

used at room temperature. Our study indicates that 7075 -T6 has higher

strength than 7475-T761, but its fatigue strength in the notched condition

is slightly lower due to its higher notch sensitivity. Alloys from the 2000

series, however, are usually preferred for cryogenic applications. Alloy

2219 -T87 is especially attractive for cryogenic aerospace applications

because it is readily weldable. This advantage is lost, however, for SSC

key collars which do not require welding.

Alloy 2090 is a relatively new material that offers a 7% lower density

and a 10% higher modulus than other aluminum alloys. Again, the advantages

are significant in aerospace applications, and not in SSC applications. At the

present time, alloy 2090 is disadvantageous in terms of availability and cost.

Service Applications

Relating laboratory data to service behavior is a practical problem

that was minimized in this study by testing a notch configuration equivalent

to that of the specific application. The square notch geometry used in this

study nearly matches the notch configuration of the service component, but

other differences must be considered. The SSC key collars are stamped, not

milled and electro-polished as our laboratory specimens were. For test

specimens, electro-polishing provides a smooth, uniform surface which

improves fatigue resistance and aids material comparisons by reducing data

scatter. Stamping may affect the performance of the manufactured parts if,

for example, residual stresses are created without subsequent stress relief

treatment. Another consideration is the manner of loading in service,

especially any difference in the fatigue (minimum- to -maximum) stress ratio;

higher stress ratio usually decrease the cycles to failure for a given range

(maximum minus minimum) stress.

138
CONCLUSIONS

The following conclusions emerge from this study of the fatigue behavior of

notched high strength aluminum alloys at 76 K:

1. In terms of yield strength, the rank in descending order is: 7075-T6

(589 MPa), 2090-T8E41 (551 MPa), 7475-T761 (549 MPa), and 2219-T87

(484 MPa)

2. In terms of notched tensile strength, the rank in descending order is:

7475-T761 (472 MPa), 7075-T6 (450 MPa), 2090-T8E41 (443 MPa), and 2219-

T87 (435 MPa).

3. In terms of notched fatigue strength for 4 x 10 4 cycles, the rank in

descending order is- 2090-T8E41 (129 MPa), 7475-T761 (109 MPa), 7075-T6

(103 MPa), and 2219-T87 (102 MPa).

4. In terms of notch-yield ratio, the rank in descending order is 2219 -T87

(0.90), 7475-T761 (0.86), 2090-T8E41 (0.80), and 7075-T6 (0.76).

5. The ratio of the notched fatigue strength at 10 s cycles to the notchedtensile strength is nearly constant (0.21-0.23).

ACKNOWLEDGEMENTS

The sheets of 2090 and 7475 were provided through the courtesy of

Dr. R. Shepard of the Alcoa Technical Center, Alcoa Center, PA. This work

was sponsored by the Superconducting Super Collider Central Design Group,

Universities Research Association, with Dr. T. Scalise as the project

monitor. We are grateful to Dr. Scalise for discussion and review the

manuscript.

SYMBOLS USED

E — Young's modulus

EL - tensile elongation

NFS - notched fatigue strength

NFSR - notched fatigue strength ratio (NFS/NTS)

NSR - notch- strength ratio (NTS/UTS)

NTS - notched tensile strength

NYR - notch-yield ratio (NTS/YS)

UTS - ultimate tensile strength

YS - yield strength

139
REFERENCES

(1) Peters, C., Mirk, K. , Wandestforde , A., and Taylor, C., presented at

the Tenth International Conference on Magnet Technology, Boston,

September 21—25, 1987.

(2) Kaufman, J.W., Sha, G.T., Kohn, R.F., and Bucci, R.J., in: Cracks and

Fracture, ASTM STP 601, Amer. Soc. Test. Maters., Philadelphia, 1976,

pp. 169-190.

(3) McCammon R.D., and Rosenberg, H.M.

,

Proc . Roy. Soc. Lond. , A, Vol. 242,

1957, pp. 203-211.

(4) Dorward, R.C., Scripta Metall., Vol. 20, 1986, pp. 1379—1383.

(5) Glazer, J., Verzasconi, S.L., Sawtell, R.R.

,

and Morris , J.W., Jr.,

Metall. Trans., V0I.I8A, 1987, pp . 1695- 1701

.

(6) Venkateswara Rao , K.T.

,

Hayashigatani , H.F., Yu, W.

,

and Ritchie, R.O.

,

Scripta Metall., Vol. 22, 1988, pp. 93-98.

140
ELASTIC CONSTANTS OF FOUR Fe-Cr-Ni-Mn ALLOYS

S.A. Kima

f H.M. Ledbettera

, and Y. Li^

aInstitute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado 80303 (U.S.A.)

^Institute of Metal Research

Academia Sinica

Shenyang (P.R. China)

ABSTRACT

Using ultrasonic methods, we measured the complete

polycrystalline elastic constants of four Fe—Cr—Ni—Mn alloys.

Combining these results with those reported previously for five

similar alloys, we focused on the alloying effects of chromium,

nickel, and manganese. The alloys contain 16 to 23 atomic percent

chromium, 7 to 28 atomic percent nickel, and 0 to 13 atomic

percent manganese. Only Mn produced regular effects, consistent

with volume changes. Both Cr and Ni produced surprising increases

in the bulk modulus and equally surprising decreases in the shear

modulus. We hypothesize that changing the bonding electrons from

d-character to s-character explains such irregularities. We

fitted the measurements to Ducastelle's model, which contains two

terms: bandstructure and repulsive energies.

1. INTRODUCTION

In cryogenic applications, austenitic stainless steels are used

extensively for storage and handling because they resist corrosion and show

relatively high strength, stability, and toughness.

Austenitic stainless steels are formed by adding f. c . c . -stabilizing

elements such as nickel or manganese to the Fe-Cr system. Nickel in Fe-Cr

not only acts as a strong austenite stabilizer, it also improves material

properties [1]. Nickel shortage during World War II forced manufacturers to

reduce or replace nickel with other elements such as manganese and nitrogen.

Many questions arose. How do these replacements affect mechanical

properties such as tensile strength, yield strength, elongation,

weldability, formability , machinability , and corrosion resistance? Many

researchers [2—4] considered these effects.

Intended for Materials Science and Engineering.

141
Elastic constant changes caused by reducing or replacing nickel with

manganese in Fe-Cr-Ni were studied recently [5,6]. Elastic constants

interrelate fundamental physical properties. Also, they are key design

parameters. For cryogenic -equipment design, for example, thermal

expansivity (0) and thermal conductivity (/c) are important variables to

consider. Both 0 - ^(7,p,Cv ,B) and k - /c(9,7,T) relate to elastic

constants. We denote 7 - Gruneisen parameter, p - mass density, Gy - specific heat, 9 - pervasive parameter, T - temperature, and B - bulk

modulus. Elastic constants also relate closely to plastic -deformation

properties

.

In this study, we determined the elastic constants of four Fe—Cr-Ni-Mn

alloys. We combined these results with our previous results for five

similar alloys. From the set of nine, we deduced alloying effects of Cr,

Ni , and Mn

.

2. MEASUREMENTS

2.1 Materials

Alloys were prepared at the Institute of Metal Research (Shenyang)

.

Table 1 shows the chemical compositions. The alloys received the following

heat treatments. Alloy Fe-16Cr-28Ni: annealed 1 h at 1253 K, quenched in

water to 993 K, and cooled in air for 16 h. Alloy Fe-23Cr-12Ni-5Mn:

annealed 1 h at 1338 K and quenched in water. Alloys Fe-21Cr-9Ni-9Mn and

Fe-21Cr-7Ni-9Mn: annealed 1 h at 1323 K and quenched in water.

2.2 Sound -velocity measurements

Room- temperature ultrasonic -sound velocities were determined by a

pulse-echo-superposition technique. Quartz piezoelectric crystals with

fundamental resonances between 3 and 9 MHz were used as transducers. X-cut

and a.c.-cut transducers were used to produce longitudinal and shear waves,

respectively. A transducer was rigidly cemented to a specimen's flat and

parallel surface using phenyl salicylate as the bonding agent. Experimental

details and the methods of obtaining the correct transit time were reported

elsewhere [7]. The ultrasonic sound velocity, v, was computed from

2i (1) V

t

*

Here, i denotes specimen length and t round-trip transit time.

142
Table 1. Chemical compositions, atomic percent.

Alloy Cr Ni Mr N C Si P S Mo V Cu A1 Ti Fe

Fe-16Cr—28Ni-OMna 16.36 28.08 0.19 <0.02 0.11 0.34 0.02 0.01 0.89 0.29 - 0.41 2.16 Balance

Fe—23Cr-12Ni-5Mnb 23.48 12.04 4.97 1.22 0.21 0.57 0.04 0.01 1.49 0.24 - - - Balance

Fe-21Cr—9Ni—9Mn 20.62 8.86 8.94 0.97 0.14 0.64 0.03 0.01 - - 0.05 0.14 0.06 Balance

Fe—21Cr—7Ni-9Mn 21 . 36 7.05 9.07 1.01 0.17 0.99 0.04 <0.01 - - 0.07 <0.06 0.06 Balance

a Also contains 0.01B

b Also contains 0.17Nb

2.3 Density measurement

Using distilled water as a standard, and Archimedes's method, we

determined the specimen's density:

P

Pv

1 - (mv/ma)

( 2 )

Here, pw - density of water, - specimen mass in water, and ma - specimen

mass in air.

3 . RESULTS

The elastic moduli are related to longitudinal, v^, and transverse, vt , ultrasonic sound velocities by

longitudinal modulus

c i - pVi 2 (3)

shear modulus

G - pvt

2 (4)

3pvt

2 (vi2 - 4 vt2

) Young modulus: E — vi

2 - vt2 (5)

bulk modulus

B - P( vi

2 ~ 4 vt2

) (6)

3

Poisson ratio: _

1 (v,» - 2vr

2) (7)

Table 2 shows room- temperature elastic constants and mass density.

143
Table 2. Elastic constants of Fe—Cr—Ni—Mn alloys corrected for C+N and Mo.

Alloy Density

(g/cm3

) Cl

(GPa)

G

(GPa)

B (GPa)

E (GPa)

V

Fe—16Cr—28Ni—Oftn 8.029 269.1 72.77 172.2 191.5 0.3149

Fe-2 3Cr—1 2Ni—5Mn 7.879 259.8 72.96 162.6 190.3 0.3048

Fe-20Cr—9Ni-9Mn 7.825 253.8 76.94 151.2 197.3 0.2826

Fe—21Cr—7Ni—9Mn 7.802 254.9 76.51 153.0 196.7 0.2858

Fe—19Cr-lNi—12Mn 7.734 253.3 74.7 153.7 192.8 0.291

Fe—19Cr—3Ni—1 3Mn 7.817 247.2 75.1 147.1 192.6 0.282

Fe-22Cr—6Ni—9Mn 7.838 252.8 75.4 152.3 194.1 0.288

Fe—22Cr—12Ni—5Mn 7.884 257.9 73.9 159.4 192.1 0.299

Fe—20Cr—9Ni-2Mn 7.860 260.5 77.7 156.9 200.1 0.287

Estimated error in the velocity measurements is d(inv) - [(d(ini)) 2 +

(d(int)

)

2

] x / 2

. Both and vt of d(inv)

density measurements, we estimate d(inp)

elastic moduli are therefore 3/2|d(inCj>) - d(inG) | “ | d(inB) |«|d(inE) |«l/2 | d(in»/) | <0 . 3%

< 0.08 % are expected. On

< 0.01%. Uncertainties in the

By linear- least- squares analysis, C - C 0 + axcr + bxj^i + cxMn + e » where e is the residual error associated with the response C, we determined

the separate effects of chromium, nickel, and manganese. Table 3 shows the

results, which reflect measured values corrected for both C+N and Mo

content, using results reported previously [8,9]. The mean square error,

s

2

, is calculated from the error sum of squares with five degrees of

freedom; s

2 for C^, G, B, E, and u is 3.89, 2.70, 11.9, 10.1, and 5.6 x

10~6

, respectively. We neglected variations in other alloying elements.

TABLE 3. Compositonal variation of elastic constants based on nine Fe-CrNi-Mn alloys: C - CQ + axcr + bxjji + cx^n ; x is in atomic

percent. CQ and a, b, c, units are GPa, except dimensionless u.

C Co a x 10 3 Co

b x 10 3 Co

c x 10 3 Co

c i 264.7 -0.7682 1.025 -3.526

G 84.11 -2.556 -3.190 -3.412

B 152.3 0.5837 4.180 -3.569

E 213.9 -2.203 -2.354 -3.495

1/ 0.2651 2.675 4.946 0.2591

144
4. DISCUSSION

A simple sphere- in-hole model (Bitter [10], Eshelby [11]) leads to

expected elastic -stiffness changes related to volume changes, dV. If we

focus on the bulk and shear moduli, then it follows that, for these

materials, [8]

dV

V

( 8 )

and

M.-2.06 dV

V

' (9)

Here, we took u - 0.290 and other parameters reported previously [8]. From

the numbers in Table 4, we expect that Mn will decrease B and G, that Cr

will decrease both by a smaller amount, and that Ni will increase both.

Among these six expectations, four are met. Contrary to expectation, Cr

increases B, and Ni decreases G. Thus, a simple sphere- in-hole model, based

on volume, fails to explain the alloy behavior. This suggests strong

electronic changes.

Considering Mn, for B and G, we find its effect similar to that found

previously for Fe—Cr—Ni alloys [13] where Mn varied up to 6 at. pet. Those

alloys contained lower N content, averaging 0.6 at. pet. Despite Mn's many

peculiar elemental properties, including a large change of this system's

Neel temperature [14], it affects elastic properties in a regular, expected

way. It is interesting that Mn contains 7 bonding electrons, close to the

alloy's number: 7.8. Thus, among all possible d-electron alloying

elements, Mn may least disrupt the alloy's electronic structure. As

discussed by Pettifor [15], Mn prefers a c.p.h. or an f.c.c. crystal

structure to a b.c.c. crystal structure.

TABLE 4. Properties of constituent atoms and one reference alloy.

V

(A3

)

B (GPa)

E (GPa)G

(GPa)

i/

Fe 11.70 166.9 212.4 82.4 0.288

Cr 11.94 161.9 279.5 115.3 0.212

Ni 10.87 186.0 220.6 84.7 0.302

Mn 12.22 124.4 197.7 76.5 0.292

Fe—18Cr—8Ni a 11.59 158.2 199.6 77.4 0.290

a Refs. 7, 12.

145
Considering Cr, we find the unexpected result that B increases and that

B and G change oppositely. Almost all elastic-constant models predict the

same sign for (l/x)(dB/B) and (l/x)(dG/G). A sign difference signals an

interatomic -bonding change that affects (in this case) dilatational modes

more than shear modes. Ledbetter and Kim [9] reported similar behavior for

Mo alloyed into Fe—Cr-Ni. Cr and Mo are isoelectronic ; both represent

b.c.c. paradigms. We expect both to disrupt strongly the f.c.c. electronic

structure (interatomic bonds), and produce unusual changes in f.c.c. -alloy

physical properties. Also, we note that Cr, despite possessing a shear

modulus 50 pet. larger than the Fe—Cr—Ni alloy, decreases the alloy's shear

modulus. Mo, with a 60 pet. larger shear modulus, also decreases the

alloy's shear modulus. Thus, rule-of -mixture models fail drastically for

these d-electron alloys. Alloying alters sharply the electronic state.

Considering Ni, except for its magnetism, a typical f.c.c. element, we

find its alloying effects surprisingly large and hard to understand,

especially the sign difference in (l/x)(dB/B) and (l/x)(dG/G). Again, this

suggests an interatomic -bonding change, in this case affecting shear modes

more than dilatational modes. Thus, electronically, Ni behaves differently

from Cr. Even though both are 3d-electron transition metals, Cr possesses

only six bonding electrons while Ni possesses ten.

The present results for Ni differ from those found in another study

[16] in Fe—21Cr—xNi—9Mn (at. pet.) where Ni varied from 4 to 11 at. pet. and

the alloys contained lower N content, averaging 0.8 at. pet. For the shear

modulus, that study found (l/x)(dG/G) - 0.861 x 10“3

, much smaller than the

present result. The reason for this difference remains under study.

Above, we described irregular responses of B and G (which reflect

dilatational and shear stiffness, respectively) to alloying. By irregular,

we mean inconsistent with responses expected from volume effects, Eqs . (8)

and (9). We can achieve some understanding of these irregular responses by

realizing that B and G reflect mainly changes in the s- electrons and delectrons, respectively. The s-electrons represent a wide, spherically

symmetrical distribution. In contrast, d-electrons represent a narrow,

localized, highly directional distribution. The shear modulus G represents

resistance to changes in interatomic -bond angles (the d-electrons). The

bulk modulus B represents resistance to changes in interatomic -bond lengths

(both s-electrons and d-electrons). For transition metals, Pettifor [17]

showed that the bulk modulus arises largely from the s -electron

contribution, a contribution enhanced strongly by the ion-core orthogonality

constraint. Pettifor emphasized also that the s-electrons figure

prominently in transition-metal cohesion; they provide a repulsive pressure

that counters the strong d-electron attraction. Thus, the d-electrons often

contribute negatively to the bulk modulus. The negative contribution of delectrons to the bulk modulus arises also in the renormalized- atom

transition-metal cohesive -energy model of Gelatt et al. [18]. For the

cohesive energy, E(r)

,

they find four components: (1) renormalization

energy, (2) s-band formation energy, (3) d-band broadening energy, and (4)

s—d hybridization energy. Combined and expressed as E(r)

,

the last two

terms show a negative curvature versus r. The bulk modulus relates to E(r)

by [19]

146
B ( 10 ) -

12irr,

a

2 E(r) 2 3E(

ar 2 £(r) 1

ar J

The renormalization and s-band terms show a positive curvature; thus they

contribute positively to B. As electrons change from d-character to s- character, we expect bulk modulus to increase and shear modulus to decrease

The present study shows that both Cr and Ni cause such changes. A previous

study [9] showed the same for Mo. It remains undetermined whether changing

the d-electron character from eg to t 2 g

[20] affects the elastic constants.

These two subbands possess a different wave -function symmetry. The eg

subband shows six lobes along <1 1 1> directions. For these f.c.c. alloys,

the Young-modulus maxima occur along <1 1 1> directions.

The opposite changes in the bulk and shear moduli cause us to focus on

the Poisson ratio

1 3B - 2G

2 3B + G

' ( 11 )

In differential form, we obtain

du 9BG f dB _

dG

]

i/ ~ (3B+G) (3B—2G) [ B G J

- 0.624 f

£B

l B

( 12 )

The factor 0.624 applies to the Fe—18Cr-8Ni reference alloy. From Eq. (12)

we see that the opposite changes in B and G combine to produce a large du/u

This reminds us of the Koster—Franz [21] view that "Poisson's ratio depends

to a much greater extent on the conditions of bonding than do the other

elastic constants." Considering the Voigt bound, Ducastelle [22] derived

the following simple relationship:

1 3 - x

2 6 - 7x

(13)

Here, x - q/p, the ratio of the bands true ture- energy and repulsive -energy

exponential coefficients, with p>q. We can extend this model to the Reuss

bound:

9 — 4x

‘'R Z

33 - 38x

(14)

Finally, we can obtain the physically more realistic Hill-model result:

87 - 34x

"H “ Z • 333 - 386x

( 15 )

147
Solving for x, we obtain

- 333i/ - 87

x "

386t/ - 34 *

Thus, the Poisson ratio leads directly to the Ducas telle -model q/p ratio,

which gives the relative contribution of bandstructure energy, U^, and

repulsive energy, Ur , to cohesion [9]:

Ducastelle's model represents almost the simplest possible equilibrium

model of a transition metal, two terms: a repulsive term and an attractive

bandstructure term. Better models contain more terms. For example, the

renormalized-atom model of Gelatt et al. contains four principal energy

terms. Positive contributions to the bulk modulus arise from two terms:

renormalization and s-band. The renormalization energy relates to the freeatom wave functions. The s-band energy is given by

AES - E0 + 3/5 Ef - E| . (18)

Here, E0 denotes the bottom of the s-band, Ef denotes Fermi energy, and E£

denotes the renormalized-atom one -electron energy. Negative contributions

to the bulk modulus arise from the combined d-band broadening and s—

d

hybridization. Neither the renormalization energy nor the d-band terms show

large curvature near the equilibrium volume. Thus, the s-band term

dominates the bulk modulus. However, for copper, for example, considering

only eq. (18) predicts a bulk modulus approximately 50% too high. Thus, we

cannot neglect the the d-band terms, which contribute negatively. The large

positive curvatures of E0 and Ef near the equilibrium volume result from an

effective compression caused by s—d hybridization. (In the 3d-row, K and Cu

possess atomic volumes of 71.4 and 11.7 A

3

, respectively!)

Table 5 shows the results of fitting Ducastelle's model to our nine

alloys, to an 18Cr—8Ni alloy [7], and to seven f.c.c. elements. Columns 2-4

show the input: v - Poisson ratio, B - bulk modulus, U - cohesive energy.

Columns 5 and 6 show the percentage bandstructure contributions to U and to

B. Columns 7—9 show various forms of the p and q coefficients. As

described by Fiirth [23], pqr 2 correlates with harmonic properties such as

the bulk modulus; and (p+q)r correlates with anharmonic properties such as

thermal expansivity, thermal -expansivity temperature derivative, bulkmodulus temperature derivative, and bulk-modulus pressure derivative.

Although the bandstructure energy dominates the cohesive energy, it

contributes only 16 to 34 percent to the bulk modulus. For our nine alloys,

we note a larger variation in (p+q)r than in pqr2

. Thus, these alloys would

show larger variations in thermal expansivity than in elastic stiffness, and

the usual elastic-stiffness/thermal-expansivity empirical interrelationships

should be less reliable.

148
TABLE 5. Ducas telle -model parameters.

Material V B

(GPa)

U

(GJ/m3

) ub/u Bb/B q/p (p+q)r pqr2

Ag 0.364 101.2 27.8 0.76 0.30 0.32 13.3 32.7

Au 0.424 173.5 35.9 0.71 0.34 0.42 14.5 43.5

Co 0.310 188.9 63.8 0.84 0.23 0.19 14.1 26.6

Cu 0.341 135.3 47.5 0.79 0.28 0.27 12.3 25.6

Fe 0.314 188.3 58.7 0.83 0.24 0.20 14.4 28.9

Ni 0.302 186.0 65.0 0.86 0.22 0.17 14.6 25.8

Pt 0.395 282.7 62.2 0.73 0.32 0.38 14.4 40.9

Fe—18Cr—

8Ni 0.290 158.2 58.6 0.89 0.19 0.12 15.8 24.3

Alloy 1 0.315 172.2 59.5 0.83 0.24 0.20 13.6 26.1

2 0.305 162.6 58.1 0.85 0.23 0.17 14.2 25.2

3 0.283 151.2 58.4 0.91 0.16 0.09 17.2 23.3

4 0.286 153.0 58.2 0.90 0.18 0.11 16.5 23.7

5 0.291 153.7 57.8 0.89 0.19 0.13 15.5 23.9

6 0.282 147.1 57.9 0.92 0.16 0.09 17.2 22.9

7 0.288 152.3 58.2 0.90 0.18 0.12 15.9 23.6

8 0.299 159.4 58.1 0.87 0.21 0.15 14.6 24.7

9 0.287 156.9 58.7 0.90 0.18 0.11 16.3 24.1

5 . CONCLUSIONS

Concerning the elastic constants of these quaternary Fe—Cr—Ni—Mn

alloys, we reached four conclusions:

1. Manganese lowers all the elastic stiffnesses and the Poisson ratio.

This agrees with expectations based on volume change.

2. Both Cr and Ni increase bulk modulus and decrease shear modulus. This

agrees with results reported previously for Mo. Such opposite alloying

effects for B and G disagree with most theoretical models. A

qualitative explanation arises by invoking a change in bonding

electrons from d-character to s-character.

149
3. The strong Ni alloying effects differ from those found previously for

lower Ni concentrations

.

4. Except for Mn, which behaved regularly, the other alloying elements, Cr

and Ni, caused unusually large changes in the Poisson ratio, an elastic

constant that responds strongly to changes in the type of interatomic

bonding.

REFERENCES

1. J.H.G. Monypenny, Stainless Iron and Steel, Wiley, New York, 1931,

p. 157.

2. R.B. Gunia and G.R. Woodrow, J. Metals, 5 (1970) 413-430.

3. B.N. Ferry and J.F. Eckel, J. Metals, 5 (1970) 99-107.

4. H.I. McHenry, in Austenitic Steels at Low Temperatures, Plenum, New

York, 1983, pp. 1-28.

5. H.M. Ledbetter, Metall. Trans. A, 11 (1980) 1067-1069.

6. H.M. Ledbetter, Mater. Sci. Eng., 29 (1977) 255-260.

7. H.M. Ledbetter, N.V. Frederick, and M.W. Austin, J. Appl. Phys . , 51

(1980) 305-309.

8. H.M. Ledbetter and M.W. Austin, Mat. Sci. Eng., 70 (1985) 143—149.

9. H.M. Ledbetter and S.A. Kim, J. Mater. Res., 3 (1988) 40-44.

10. F. Bitter, Phys. Rev., 37 (1931) 1527-1547.

11. J.D. Eshelby, J. Appl. Phys., 25 (1954) 255-261.

12. H.M. Ledbetter and M.W. Austin, Metal Sci., 3 (1987) 101-104.

13. H.M. Ledbetter, J. Mater. Sci., 20 (1985) 2923-2929.

14. E.R. Jones, T. Datta, C. Almasan, D. Edwards, and H.M. Ledbetter,

Mater. Sci. Eng., 91 (1987) 181—188.

15. D.G. Pettifor, J. Phys. C, 3 (1970) 367-377.

16. H.M. Ledbetter, Ming Lei, and Yiyi Li, to be published.

17. D.G. Pettifor, J. Phys. F, 8 (1978) 219-230.

18. C.D. Gelatt, H. Ehrenreich, and R.E. Watson, Phys. Rev. B, 15 (1977)

1613-1627.

19. S. Raimes, The Wave Mechanics of Electrons in Metals, North-Holland,

Amsterdam, 1961, p. 275, eqn. 9.90.

150
20. J.B. Goodenough, Phys . Rev., 120 (1960) 67—83.

21. W. Koster and H. Franz, Metall. Rev., 6 (1961) 1-55.

22. F. Ducastelle, J. Physique, 31 (1970) 1055-1062.

23. R. Furth, Proc. Roy. Soc., A183 (1944) 87-110.

151
.
NICKEL EFFECT ON ELASTIC CONSTANTS OF Fe-Cr-Ni ALLOYS

M. Lei and H. Ledbetter

Institute for Materials Science and Engineering

National Institute of Standards and Technology

Boulder, Colorado 80303, USA

Y. Li

Institute of Metal Research

Shenyang, People's Republic of China

Using a kilohertz -frequency standing-wave method, we

determined the effect of Ni content on the elastic constants of

Fe—Cr-Ni alloys containing nominally 20.3 mass percent Cr. Nickel

varied from 6.1 to 11.8 mass percent. As expected, we found

little change. For example, for five alloys, the Young modulus

averaged 196.1 ± 1.7 GPa. Similarly, the shear modulus averaged

77.3 ± 0.4 GPa. Internal friction measured both in extension and

torsion showed no systematic composition dependence. The small

elastic-constant changes agree with deductions based on atomic

size. However, the small decrease in bulk modulus and

accompanying increase in shear modulus deserve focus , which we

provide by using a Ducastelle model for transition-metal alloys.

The model contains two terms: bandstructure and ion-corerepulsion.

Effects of carbon and nitrogen [1] and manganese [2] on elastic

properties of Fe—Cr—Ni alloys were reported. Also, one of the authors [3]

reported measurements and modeling of molybdenum- alloying effect on

Fe—Cr—Ni- alloy elastic constants. Here we report new measurements and

modeling on the effect of nickel.

Materials were produced at the Institute of Metal Research, Shenyang,

China. Table 1 shows their chemical compositions.

Intended for publication in Materials Science and Engineering.

153
Table 1. Chemical composition (mass %), balance Fe.

Alloy Cr Ni Mn C Si S P N

21 20.02 6.13 9.29 0.020 0.51 0.011 0.011 0.15

22 20.15 6.20 9.35 0.035 0.58 0.010 0.013 0.22

24 20.25 8.02 9.25 0.017 0.51 0.013 0.009 0.20

25 20.43 10.17 9.00 0.014 0.48 0.013 0.009 0.20

29 20.37 11.77 9.22 0.018 0.40 0.015 0.010 0.28

We determined the Young and shear moduli and the associated internal

friction by using a Marx kilohertz- frequency resonance method [4, 5].

Measurement details were described elsewhere [6]

.

Here we give expressions

for Young modulus, E, and torsional modulus (shear modulus), G:

E - 4pf§l§ - 4pll[f0

s + (f0* - f

q

2

) 5a], (1)

and

G - 4pf|l| - 4pl| [fo 2 + (fo 2 ~ fq 2)

sf (^f) ’]

• (2)

Here, p and 1 denote mass density and specimen length; f denotes frequency;

m and d indicate mass and diameter. The subscripts s, o, and q denote

specimen, complete oscillator, and quartz crystal. The specimen's internal

friction, Qs_1

, is determined by

Qs"

1 - Qo"1 + (Qo"”1 " Qq_1 ) m

q

/ms

. (3)

Table 2 shows the principal measurement results. Since the materials

are quasi- isotropic , we invoke the well-known relationships for bulk

modulus, B, and Poisson ratio, u:

B

GE

3 (3G - E) ’ (4)

and

Figure 1 shows elastic moduli and Poisson's ratio versus nickel content. We

see that nickel increases slightly the Young and shear moduli and decreases

both bulk modulus and Poisson ratio. To verify that the changes are real,

not measurement error, we focused on measurement inaccuracy. For the

longitudinal mode, we repeated one measurement ten times and found that the

standard deviation is 0.08 GPa. For the shear mode, we repeated one

measurement six times and found a standard deviation of 0.05 GPa.

154
Table 2. Effect of nickel on Fe-Cr—Ni- alloy elastic constants and internal

friction.

Alloy Ni

(mass pet.)

P

(g/cm3

)

E

(GPa)

G

(GPa)

Qe*

(i<rs

) Qg 1 (10_s

)

21 6.13 7.786 195.0 76.9 2.9 5.7

22 6.20 7.681 195.0 10.5

24 8.02 7.764 195.6 77.4 5.5 5.4

25 10.17 7.772 195.8 5.9

29 11.77 7.759 195.6 77.6 6.9 2.6

Studies by Ledbetter and Austin [7] showed that nickel slightly

decreases the effective atomic volume. Volume changes usually cause

physical -property changes. For example, consider the bulk modulus of Fe.

From 0 to 300 K, it decreases 2.95%, while the volume increases 0.59%.

Ignoring the thermal contribution to elastic softening, we obtain

AB/B _

0.0295

AV/V " 0.0059 ( 6 )

Thus, we expect the bulk modulus to increase approximately 5% per percent

volume decrease.

A quantum -mechanical theory by Ducastelle [8] leads to simple

relationships among volume, bulk modulus, and cohesive energy. In the

theory, the total energy, U, arises from two sources: d-band energy, Ub , calculated by a tight -binding approximation and Bom-Mayer ion- core

repulsive energy, Ur

. Thus,

U - Ub + Ur - —

W

0 e”

Ce~Pr . (7)

Here , W0 , C , p , and q are constants ; and r denotes the atomic radius . A

simple expression arises for the equilibrium radius:

( 8 )

For the bulk modulus, Ducastelle obtained

B Br + Bb

pqr 0 ur. 9V0 ( p q

)

(2 + pr0 - 2 - qr 0 ) (9)

Here, Uc denotes cohesive energy, and V0 - 4jrr 0

3 /3 denotes atomic volume.

Thus, Ducastelle' s model predicts that decreased atomic volume leads to

increased bulk modulus. If volume changes little, then we must consider

changes in other parameters

155
200

NICKEL (mass %)

Fig. 1. Compositional variation of E - Young modulus, B - bulk modulus,

G - shear modulus, and u - Poisson ratio.

156
20

CM

Fig. 2. Compositional variation of pr 0 /7v0 /Uc and qr 0 /7v0 /Uc

.

157
Despite some limitations of Ducastelle's model, we can use it to

estimate pr0 and qr0 , parameters characterizing the repulsive and bandstructure energies. For this estimate, besides Eq. (9), we need another

relationship among p, q, and physical properties. We choose the Poisson

ratio v. Using Ducastelle's expression for u:

and

qr 0

6 - 7 (q/p)

9

Eq. (9), we <

fl4i/ - 1 BV

L12*/ - 3 u(

fl2i/ - 3 BVp

L14i/ - 1 uc

-

( 10 )

(ID

( 12 )

In this study, we ignore the volume change caused by nickel alloying and

take V0 /Uc as a constant. Figure 2 shows that p increases and q decreases

as nickel increases. Since p and q characterize repulsive and bandstructure

energy, respectively, added nickel to Fe—Cr—Ni alloy will decrease the

contribution of repulsive energy and increase that of the bandstructure

energy to the total energy.

In summary, we find that Ni alloyed into f.c.c. Fe—Cr—Ni alloys

produces small elastic-constants changes, consistent with small atomicvolume change. However, the small bulk-modulus decrease coupled with a

small shear-modulus increase provides a surprise and a challenge to theory.

Applying Ducastelle's model suggests that adding nickel increases the

bandstructure: core -repuls ion energy balance.

ACKNOWLEDGMENT

M. Lei is a visiting scientist from the Institute of Metal Research

(Shenyang)

.

REFERENCES

1. H.M. Ledbetter and M.W. Austin, Mater. Sci. and Eng., 70 (1985)

143-149.

2. H.M. Ledbetter and S.A. Kim, J. Mater. Sci., 21 (1988) 2129-2132.

3. H.M. Ledbetter and S.A. Kim, J. Mater. Res., 3 (1988) 40—44.

4. J. Marx, Rev. Sci. Instrum., 22 (1951) 503-509.

5. W.H. Robinson and A. Edgar, IEEE Trans. Sonics Ultrason. , SU-21 (1974)

98-105.

6. H.M. Ledbetter, Ming Lei, and M.W. Austin, J. Appl. Phys., 5j? (1986)

1972-1976.

7. H.M. Ledbetter and M.W. Austin, National Institute of Standards and

Technology, Boulder, Colorado (1986), unpublished.

8 . F. Ducastelle, J. Phys., 2i (1970) 1055—1062.
NITROGEN EFFECT ON ELASTIC CONSTANTS OF F.C.C. Fe-18Cr-19Mn ALLOYS

S. Lin and H. Ledbetter

Institute for Materials Science and Engineering

National Institute of Standards and Technology

Boulder, Colorado 80303, U.S.A.

ABSTRACT

Previously, we studied effects of interstitial carbon-plus - nitrogen on the elastic constants of f.c.c. Fe-18Cr-10Ni-lMn

alloys. Consistent with a volume increase, all the elastic

stiffnesses decrease with increasing C+N. The present alloys show

different behavior: even though volume increases, interstitial

nitrogen atoms increase the bulk modulus . We describe that 3delectron elements show a peculiar bulk-modulus-electronconcentration behavior (B versus ne ) . At first B increases with

increasing r^; beyond a critical concentration, B decreases

rapidly. Applying Ducastelle's model (bands true ture and repulsion

energies) shows that interstitial nitrogen increases the

bandstructure contribution to the bulk modulus.

1 . INTRODUCTION

Adding f.c.c. stabilizing elements such as nickel or manganese to Fe-Cr

alloys produces various types of austenitic stainless steels, which show

different physical properties. The interstitial atoms—nitrogen and

carbon—stabilize the f.c.c. lattice and improve the strength, toughness,

and corrosion resistance [1]. Nitrogen- strengthened austenitic stainless

steels find extensive low- temperature applications, and they interest many

researchers [1—5].

Ledbetter and coworkers [5] reported the effects of carbon and nitrogen

interstitials on the elastic constants of AISI-304 stainless steels

(Fe—18Cr—lONi-IMn) . They found that carbon and nitrogen decrease all the

elastic stiffnesses.

In this study, we considered higher-Mn alloys and found big differences

between Fe—18Cr—19Mn and Fe-18Cr-10Ni-lMn. Especially, nitrogen

interstitial atoms affect the elastic constants differently. We think that

the lower Ni content produces a negligible elastic-constant change.

Intended for Materials Science and Engineering .

159
TABLE 1

Chemical composition of high-Mn stainless steels (at.%), balance Fe

Alloy N C Si Mn Cr Ni Mo V

1 1.51 0.06 0.638 19.97 14.62 0.296 0.079 0.117

2 2.22 0.62 0.668 18.24 18.96 0.210 0.045 0.053

3 3.05 0.18 0.495 19.85 19.22 0.528 0.006 0.126

4 3.70 0.27 0.991 18.89 18.01 0.118 1.160 0.147

2 . MEASUREMENTS

2.1 Materials

We obtained four stainless steels from a commercial source. They had a similar base composition, but different interstitial-atom content. Table 1 shows the detailed chemical compositions.

2.2 Sound -velocity methods

We used a pulse -echo -overlap technique [6] to measure the longitudinal

and transverse sound velocities at ambient temperature. The specimens were

cut into 1.5 -cm cubes and were ground so that opposite faces were flat and

parallel within 5 jim. Quartz piezoelectric crystals with fundamental

resonances between 2 and 7 MHz were glued to the specimens with phenyl

salicylate. An x-cut transducer was used for longitudinal waves and an accut transducer for transverse waves. Further experimental details were

given elsewhere [6]. The sound velocity v was obtained from

v - 2i/t . (1)

Here, i denotes the specimen length and t the round-trip transit time.

2.3

Mass-density measurement

We used Archimedes's method to determine the specimen's mass density.

Details are described elsewhere [7].

3 . RESULTS

Elastic - stiffness constants were derived from the general relationship

Ci - pVi 2

. (2)

160
Here, p denotes mass density and subscript i denotes the mode of the sound

wave. For the usual quasi isotropic elastic constants we have the following

relationships:

longitudinal modulus - - Pvi

2

, (3)

shear modulus - G - pvt

2

, (4)

bulk modulus - B - — (4/3)G, (5)

Young's modulus - E - 3GB/(Cjj—G) , (6)

Poisson's ratio - u - (E/2G)—

- h(C i-2G)/(C i-G). (7)

Here, and vt denote the longitudinal and transverse sound velocities.

Table 2 shows the results for p, Cj>, G, B, E, and u

Figure 1 shows

the relative elastic -constant changes with carbon-plus -nitrogen content.

Table 3 gives the least- squares results obtained by a linear fit:

Table 2. Mass density and elastic constants at T - 295 K

Alloy N+C

(at.%)

Density

(g/cm3

) vi

(km/s)

vt

(km/s)

C i

(GPa)

G

(GPa)

B

(GPa)

E

(GPa)

i/

1 1.575 7.765 5.607 3.188 244.1 78.92 138.9 199.1 0.261

2 2.843 7.716 5.655 3.172 246.8 77.64 143.2 197.3 0.270

3 3.230 7.690 5.671 3.163 247.3 76.94 144.7 196.1 0.274

4 3.972 7.671 5.697 3.146 249.0 75.92 147.7 194.5 0.281

INTERSTITIAL CONTENT (at.%)

Fig. 1. Variation of elastic constants with interstitial carbon-plus

nitrogen content for high-manganese austenitic stainless steels at

295 K.

161
TABLE 3. Least- squares parameters for various elastic constants at 295 K

ci

(GPa)

G

(GPa)

B

(GPa)

E

(GPa)

V

C(0) 241.0 80.98 133.0 202.3 0.248

a 2.005 -1.247 3.667 -1.912 0.00811

100 • a/C(0) 0.832 -1 . 540 2.758 -0.945 3.269

C - C 0 + axc+N . (8)

Here, xc+jj denotes atomic percentage of interstitial atoms: carbon plus

nitrogen. Table 3 also shows 100-a/C 0 , the percentage change in elasticconstant per percentage change in alloying element.

To calculate the unit-cell size we used the following formula:

a

3 - V - ^ . (9)

P

Here, a denotes the unit-cell dimension, V the related volume, p the mass

density, M the effective atomic mass, N the effective number of atoms in each

unit cell.

Figure 2 shows the unit-cell size versus combined interstitial-atom

content

.

4. DISCUSSION

High -manganese austenitic stainless steels (Fe—18Cr—19Mn) show different

interstitial -alloying effects from low-manganese austenitic stainless steels

(Fe—18Cr—lONi—lMn)

.

As in Fe-18Cr-10Ni-lMn, in Fe-18Cr-19Mn the interstitial carbon-plus

nitrogen increases the volume. In Fig. 2, two lines show the linear- least

squares analysis results for these two types of alloys; the solid line

corresponds to Fe—18Cr—19Mn; the dashed line corresponds to Fe—18Cr-10Ni-lMn

[8]. For these two cases, these results indicate the same interstitial

induced volume change. In Fe—18Cr—19Mn alloys, the interstitial carbon-plus

nitrogen decreases the shear modulus and Young's modulus but increases the

bulk modulus. In Fe-18Cr-10Ni-lMn alloys, interstitial carbon plus nitrogen

decreases all the elastic stiffnesses. Eshelby's model based on volume change

explains these decreases.

Small volume changes usually cause large physical -property changes. In

unalloyed iron, pressure studies [9] show that

29 (10)

AV/V AP

'

162
INTERSTITIAL CONTENT (at.%)

Fig. 2. Unit-cell size versus combined interstitial-atom content. Dashed

line indicates low-manganese austenitic stainless steels. Filled

symbols from x-ray diffraction. Open symbols from macroscopic mass

density.

Going from alloy 1 to alloy 4, the volume change is

AV/V - 3Aa/a - 2.2%. (11)

If we use Eq. (10) to estimate the interstitial-atom- induced volume

effect, we get

AB/B 5.29 AV/V 12%. (12)

This contrasts with the measured value: 6.3%, with opposite sign! The simple

volume effect fails to explain this unusual bulk-modulus behavior.

By measurement, we obtained (1/xjq+c) (AB/B) - 0.028, where xjj+c denotes

interstitial concentration. Compared with the low-manganese case where

(

/xjj+c) (AB/B) - -0.0085 or an Eshelby-model prediction ( 1/xfl+c) (AB/B) - —0.0104, this is too high in absolute value by a factor of approximately 3,

and is opposite in sign. For the shear modulus, we have (l/xjj+c ) (AG/G) - —0.015, which is almost two times larger than —0.0052 found in low-manganese

stainless steels. But it comes close to an Eshelby-model prediction:

(1/xjj+c) (AG/G) - -0.012.

163
For the zero- interstitial -content elastic constants, we get B 0 - 133.0

GPa, G 0 — 81.0 GPa, E0 — 202.3 GPa, and u 0 - 0.248. For the low-manganese

case, we get B0 - 159.6, G 0 - 78.8, E0 - 203.2, and u 0 - 0.280. Thus, Mn doe

not change much the shear and Young's moduli, but it decreases the bulk

modulus 17%. This large change correlates with manganese's large atomic

volume, which expands the lattice and makes the lattice more compressible. A

low Poisson ratio suggests weak interatomic bonding caused by increased atomi<

separation, as shown in Fig. 2.

In these alloys, the interstitials strongly affect the Poisson ratio.

Measurement shows (1/xfj+c )(A*//i/) - 0.0327, compared to -0.0015 for lowmanganese stainless steels, a factor of approximately —22! In Fe—18Cr—

lONi-IMn alloys [10], when an interstitial C or N atom is added to the f.c.c.

unit cell, the compound Fe 4 X or FeX exists locally (where X denotes C or N

interstitial atom) . The existence of intermetallic compounds means localized

electrons. Localized-electron bonding lowers Poisson's ratio and increases

the shear modulus but not the bulk modulus. In high -manganese alloys, this

kind of compound may exist, but the bulk-modulus increase and the shearmodulus decrease indicate weaker covalent bonding or less- localized electrons.

The large Poisson' s-ratio increase shows a change in atomic bonding [11]. We

suppose that in high -manganese stainless steels M-X electrons behave more

itinerant than local. In coherent-potential-approximation calculations, Ohta

[12] thinks that carbon and nitrogen simply provide extra valence electrons.

These additional itinerant electrons would increase the bulk modulus.

Cho [13] showed a linear correlation between bulk modulus B and netbonding- valence -electron concentration ne

. For nonmagnetic 3d-electron

elements, a linear correlation exists; but for magnetic 3d-electron elements,

B first increases linearly with increasing ne

. Beyond a critical

concentration near 5.3 x 10 29 m”3

, B decreases rapidly. We calculated the netbonding-electron concentration of the present four alloys, and we found that,

going from alloy 1 to alloy 4, ne changes from 5.46 x 10 29 m

—

3

to 5.39 x

10 29 m“3

, while the bulk modulus changes from 138.9 to 147.4 GPa. These alloys

behave similar to 3d-electron elements. Figure 3 shows the correlation

between B and ne for 3d-electron elements, Fe—18Cr—19Mn alloys, and

Fe—18Cr—lONi—lMn alloys. To make these calculations, we adopted Cho's

suggested valences. The free-electron model fails to explain the irregular

behavior of 3d-electron elements or alloys. Such a simple model usually

applies only to sp-electron metals. We note that ne decreases with increasing

nitrogen. The extra electrons are more than offset by the increasing volume.

Also, a small electron-concentration change causes a large bulk-modulus

change

.

Now, we consider another viewpoint: electrons in a simple band model.

Ducastelle's [15] model of cohesion uses a simple interatomic potential:

U - Ub + Ur Woe”^ + Ce"Pr . (13)

Here, U denotes total energy, Ub d-band energy, Ur repulsive energy, r interatomic spacing, and W0 , C, p, q constants.

164
Fig. 3. Correlation between bulk modulus B and net-bonding-valence-electron

concentration for 3d-electron metals and alloys. The 7-Fe point

came from reference [14]. By magnetic we mean that the material

shows either ferromagnetism or antiferromagnetism at ambient, or

lower, temperatures. The nonmagnetic—magnetic boundary represents a

rough approximation.

Considering the Voigt bound, Ducas telle obtained for the Poisson ratio

1 3—

2

* 6—7x

(14)

where x - q/p. Solving for x, we get

12t/-3

14i/—

( 15 )

165
Then the Poisson rstio leads directly to the Ducastelle -model ratio q/p, which

shows the relative contribution of bandstructure energy and repulsive

energy Ur to cohesive energy [16]:

In Ducastelle's model

B - V0 pqr<> 2u« 9V0

(17)

Here, V denotes volume and subscript o represents the equilibrium state. And

we can separate p and q by using Eqs . (15) and (17).

Table 4 shows the results of fitting Ducastelle's model to our four

alloys. Columns 5 and 6 show the percentage bandstructure contributions to U

and B. Columns 7—9 show several forms of the p and q coefficients, where pqr 2 correlates with harmonic properties such as the bulk modulus, (p+q)r

correlates with anharmonic properties such as thermal expansivity, thermal

expansivity temperature derivative, bulk-modulus temperature derivative, and

bulk-modulus pressure derivative [17].

Although the bandstructure energy dominates the cohesive energy, it

contributes only 12 to 19 percent to the bulk modulus. Larger variation in

(p+q)r than in pqr2 suggests a arger variation in thermal expansivity than in

elastic stiffness.

With increasing interstitial content, bandstructure energy contributes

less to cohesion, but contributes more to the bulk modulus, contrary to many

simple models where bulk modulus is proportional to cohesive energy.

TABLE 4. Ducas telle -model parameters

Alloy i/ B

(GPa)

Uo/Vo

(GJ/ms

)

Ub/U Bb/B q/p (p+q)r pqr 2

1 0.261 138.9 57.9 0.95 0.12 0.05 21.9 21.6

2 0.270 143.2 57.4 0.92 0.16 0.09 17.5 22.5

3 0.274 144.7 57.3 0.91 0.17 0.10 16.5 22.7

4 0.281 147.4 57.0 0.89 0.19 0.13 15.3 23 .

3

166
5. CONCLUSIONS

From this study, we reached four principal conclusions:

1. Alloying interstitial carbon-plus -nitrogen into high -manganese Fe—18Cr-19Mn

alloys decreases the shear modulus and Young's modulus, but increases bulk

modulus, which is inconsistent with the volume increase. The shear-modulus

change agrees with Eshelby's model. In a previous study on the effects of

interstitial carbon-plus -nitrogen in Fe—18Cr-10Ni-lMn alloys, all the elastic

stiffnesses decreased with increasing C+N, consistent with volume effects.

2. Interstitial carbon-plus -nitrogen increases Poisson's ratio remarkably; this

indicates interatomic -bonding change. The interstitial -contributed bonding

electrons behave itinerantly, increasing the bulk modulus.

3. The 3d-electron elements show a peculiar bulk-modulus—electron-concentration

behavior. At first B increases with increasing ne ; beyond a critical

concentration, B decreases rapidly. The present alloys contain higher electron

concentration than the critical, and they behave like the magnetic 3d- electron

elements

4. Applying Ducastelle's model (repulsion and bandstructure energies) suggests tha

interstitial carbon-plus -nitrogen decreases the contribution of bandstructure

energy to cohesion and increases the contribution to the bulk modulus.

ACKNOWLEDGMENTS

Our thanks to M.W. Austin and S.A. Kim, who assisted with measurements. Sihan Lin i a visiting scientist from Tsinghua University, Beijing, P.R. China. Alloys for the study

came from M. Harzenmoser and P. Uggowitzer (Zurich). Partial support for this study came

from the DoE Office of Fusion Energy.

REFERENCES

[1] P.J. Uggowitzer and M. Harzenmoser, in High Nitrogen Steels, The Institute of Metals

London, 1989, p. 174.

[2] R.P. Reed and N.J. Simon, in High Nitrogen Steels, The Institute of Metals, London.

1989, p. 180.

[3] K. Fritscher, in High Nitrogen Steels, The Institute of Metals, London, 1989 p 208

[4] H.M. Ledbetter, Mater. Sci. Eng., 29 (1977) 255-260.

[5] H.M. Ledbetter and M.W. Austin, Mater. Sci. Eng., 70 (1985) 143-149.

[6] H.M. Ledbetter, N.V. Frederick, and M.W. Austin, J. Appl. Phys., 51 (1980) 305-209

[7] S.A. Kim, H.M. Ledbetter, and Yiyi Li, unpublished research (1989).

[8] H.M. Ledbetter and M.W. Austin, Mater. Sci. Tech., 3 (1987) 101—104.

167
[9]

M.W. Guinan and D.N. Beshers, J. Phys. Chem. Solids, 29 (1968) 541-549.

[10] H.M. Ledbetter, M.W. Austin, and S.A. Kim, Mater. Sci. Eng., 85 (1987) 85-89.

[11] W. Kdster and H. Franz, Metall. Rev., 6 (1961) 1-55.

[12] Y. Ohta (Nagoya University), personal communication (1989).

[13] S.-A. Cho, Trans. Jap. Inst. Met., 22 (1981) 643-652.

[14] H.M. Ledbetter, Phys. Stat. Solidi (a), 85 (1984) 89-96.

[15] F. Ducas telle , J. Physique, 31 (1970) 1055-1062.

[16] H.M. Ledbetter and S.A. Kim, J. Mater. Res., 3 (1988) 40—44.

[17] R. Fiirth, Proc. Roy. Soc., A183 (1944) 87—110.

168
LOW -TEMPERATURE MAGNETIC -ELASTIC ANOMALIES

IN F.C.C. Fe—Cr—Ni ALLOYS

H. LEDBETTER

Institute for Materials Science and Engineering

National Institute of Standards and TechnologyBoulder, Colorado, 80303, USA

Using ultrasonic methods, we measured the complete elastic

constants of several polycrystalline face-centered-cubic Fe—Cr—Ni

alloys at low temperatures: 4-295 K. The alloys represent nine

interstitial (carbon-plus -nitrogen) compositions. We also measured d.c. magnetic susceptibility. Upon cooling, a paramagnetic

antiferromagnetic (Neel) phase transition occurs. The Neel

transition temperature depends strongly on interstitial content.

We fit both composition and temperature effects to Ducastelle's

3d-electron model, which contains two terms: bandstructure and

repulsive energies.

1 . Introduction

The study reported here represents part of a larger study to determine

and understand the low- temperature elastic -magnetic properties of f.c.c.

Fe—Cr—Ni alloys, which provide the basis for commercial austenitic steels.

For the same nine alloys, which contain 0.08 to 0.36 mass percent (0.34 to

1.47 atomic percent) of interstitial carbon-plus -nitrogen, there exist

several previous reports: ambient -temperature elastic constants [1],

liquid-helium- temperature elastic constants [2], low- temperature magnetic

susceptibility [3], and low- temperature neutron diffraction [4].

Here, we analyze the elastic constants by Ducastelle's model [5].

Considering the substitutional element, the studied alloys [1] consisted on

average of Fe—19 .4Cr-9 . 5Ni-l . 6Mn (atomic percent). These were nine carbonplus -nitrogen contents: 0.34, 0.45, 0.69, 0.79, 0.88, 0.95, 1.29, 1.38,

1.47 (atomic percent). For this Fe—Cr—Ni composition, the alloys are

unstable at low temperatures. The interstitial C+N stabilizes the f.c.c.

crystal structure. We found no evidence of crystal -structure transitions.

Intended for publication in Physica B; proceedings of ISOMES '89.

169
2. Elastic and magnetic results

Figure 1 shows, for alloy 1, the temperature variation of four polycrystalline quasi isotropic elastic constants. Upon cooling the alloy, at

the N4el temperature (47.9 K) , all the elastic stiffnesses (dilatational and

shear) soften. Below Tjj, the stiffnesses, especially the bulk modulus, tend

to resume a normal temperature dependence. Above Tfj, the Poisson ratio

behaves normally; but, below it shows an abnormal slope indicating

continued magnetic changes. Also, the bulk modulus, representing dilatational stiffness, begins to soften well above T^, at approximately 125 K.

The heating and cooling curves overlap, indicating zero hysteresis. Near

T^, using x-ray diffraction, we detected neither a volume change nor a

change from cubic symmetry [6].

1.10

1.08

1.06 -

1.04

A

.aDoog0

G

Q°°Or>A

^ DQn

A

A °Oq£a

C (T) 1 02 - C ( 29S)

1.00 Ooooooooq.£

00000°000<^

>00000

000000000000o0000o°°oo^g0.98 - w

VVv*

V*wVVV

r

0.94 1 L

100 200

TEMPERATURE (K)

300

Fig. 1. For alloy 1, temperature variation of four polycrystalline elastic

constants: G - shear modulus, E - Young modulus, B - bulk

modulus, v - Poisson ratio. Values normalized to ambient temperature. Broken vertical lines indicate Neel temperature.

170
Fig. 2. For alloys 1, 5, and 8, temperature variation of d.c. magnetic

susceptibility. Hagnetic transition temperature shifts systematically with increasing interstitial -atom content. Specimens were

cooled to 5 K in zero field and measured during warming in 1.0 mT.

Figure 2 shows typical magnetic -susceptibility^temperature curves. All

these alloys show a distinct peak as a function of temperature. The

behavior characterizes a paramagnetic- to -antiferromagnetic phase transition

[7]. Above the transition temperature, the susceptibility is well described

by a Curie -Weiss law:

X(T) - Xo + . (1)

Here, Xo » C, and © denote constants, which we reported previously [3].

The temperature- independent xo results from underlying diamagnetism and

Pauli spin paramagnetism caused by the lattice dilation and the lower Fermi

energy. The average effective moment per atom obtained from the average

value of C was 1 . 14 ± 0.07 jig. Neutron diffraction [4] gave a lower moment: * 0.5 This indicates weak external -electron s—d hybridization.

3. Ducas telle model

Ducastelle [5] considered a simple interatomic potential:

U - Ub + Ur - -Ae“^r + De”Pr

. (2)

171
Here, U denotes energy, subscript b bands trueture , subscript r repulsive,

and r interatomic spacing. Considering the Voigt bound, Ducastelle derived

the following simple relationship for the Poisson ratio:

1 3 - x

I'V “ • (3)

2 6 — 7x

Here, x - q/p, the ratio of the bands true ture- energy and repulsive -energy

exponential coefficients, with p>q. We can extend this model to the Reuss

bound

9 — 4x

~ 33 - 38x (4)

Finally, we can obtain the physically more realistic Hill-model result:

87 - 34x

^H " 333 _ 386x (5)

Solving for x, we obtain

333i/ - 87

386i/ - 34

‘ ( 6 )

Thus, the Poisson ratio leads directly to the Ducastelle -model q/p ratio,

which gives the relative contribution of bandstructure energy U^, and

repulsive energy, Ur , to cohesion:

ur

(7)

We can separate p and q by considering the bulk modulus. In Ducastelle'

s

model

B - V0 a

3 U~l rn

'2 fanjl pqr„ 2Un av2Jv_v0 " 9V0 [ar 2 J r-r0 " 9V0

( 8 )

Here, V denotes volume and subscript o the equilibrium state.

4. Model- fitting results and discussion

Table I shows the results of fitting Ducastelle 's model to our nine

alloys, to an 18Cr—8Ni alloy [8], and to seven f.c.c. elements. Columns 2—4

show the input: u - Poisson ratio, B - bulk modulus, U - cohesive energy.

Columns 5 and 6 show the percentage bandstructure contributions to U and to

B. Columns 7—9 show various forms of the p and q coefficients. As des172
cribed by Forth [9]

,

pqr 2 correlates with harmonic properties such as the

bulk modulus; and (p+q)r correlates with anharmonic properties such as

thermal expansivity, thermal -expansivity temperature derivative, bulkmodulus temperature derivative, and bulk-modulus pressure derivative.

Although the bands tructure energy dominates the cohesive energy, it

contributes only 16 to 34 percent to the bulk modulus. For our nine alloys,

we note a larger variation in (p+q)r than in pqr 2

. Thus, these alloys would

show larger variations in thermal expansivity than in elastic stiffness, and

the usual elastic-stiffness—thermal-expansivity empirical interrelationships

should be less reliable.

TABLE I. Ducas telle -model input and derived parameters •

Material V B

(GPa)

U

(GJ/m3

) ub/u Bb/B q/p (p+q)r pqr 2

Ag 0.364 101.2 27.8 0.76 0.30 0.32 13.3 32.7

Au 0.424 173.5 35.9 0.71 0.34 0.42 14.5 43.5

Co 0.310 188.9 63.8 0.84 0.23 0.19 14.1 26.6

Cu 0.341 135.3 47.5 0.79 0.28 0.27 12.3 25.6

Fe 0.314 188.3 58.7 0.83 0.24 0.20 14.4 28.9

Ni 0.302 186.0 65.0 0.86 0.22 0.17 14.6 25.8

Pt 0.395 282.7 62.2 0.73 0.32 0.38 14.4 40.9

Fe—18Cr—

8Ni 0.290 158.2 58.6 0.89 0.19 0.12 15.8 24.3

Alloy 1 0.315 172.2 59.5 0.83 0.24 0.20 13.6 26.1

2 0.305 162.6 58.1 0.85 0.23 0.17 14.2 25.2

3 0.283 151.2 58.4 0.91 0.16 0.09 17.2 23.3

4 0.286 153.0 58.2 0.90 0.18 0.11 16.5 23.7

5 0.291 153.7 57.8 0.89 0.19 0.13 15.5 23.9

6 0.282 147.1 57.9 0.92 0.16 0.09 17.2 22.9

7 0.288 152.3 58.2 0.90 0.18 0.12 15.9 23.6

8 0.299 159.4 58.1 0.87 0.21 0.15 14.6 24.7

9 0.287 156.9 58.7 0.90 0.18 0.11 16.3 24.1

173
Fig. 3. Compositional variation of bandstructure contribution to bulk

modulus. The low- temperature (antiferromagnetic) variation

differs sharply from the ambient- temperature (paramagnetic)

variation.

Ducastelle's model represents almost the simplest possible equilibrium

model of a transition metal, two terms: a repulsive term and an attractive

bandstructure term. Better models contain more terms. For example, the

renormalized- atom model of Gelatt and coworkers [10] contains four principal

energy terms. Positive contributions to the bulk modulus arise from two

terms: renormalization and s-band. The renormalization energy relates to

the free-atom wave functions. The s-band energy is given by

AES - E„ + (3/5) Ef - El . (9)

Here, E0 denotes the bottom of the s-band, Ef denotes Fermi Energy, and E|

denotes the renormalized- atom one -electron energy. Negative contributions

to the bulk modulus arise from the combined d-band broadening and s-d

hybridization. Neither the renormalization energy nor the d-band terms show

large curvature near the equilibrium volume. Thus, the s-band term dominates the bulk modulus. However, for copper, for example, considering only

eq. (9) predicts a bulk modulus approximately 50% too high. Thus, we cannot

neglect the d-band terms, which contribute negatively. The large positive

curvatures of E0 and Ef near the equilibrium volume result from an effective

compression caused by s-d hybridization. (In the 3d-row, K and Cu possess

atomic volumes of 71.4 and 11.7 A

3

, respectively!)

Figure 3 shows the composition variation of the bandstructure component

of the bulk modulus. At low temperatures, contributes less to B.

Surprisingly, the 295-K and 4-K variations differ strongly, indicating

different alloying (perhaps volume) effects at the two temperatures.

174
Fig. 4. For alloy 1, temperature variation of Ducas telle -model parameters.

Neel temperature equals 47.9 K.

Figure 4 shows the temperature variations of the Ducas telle -model

parameters. As temperature decreases, p increases while q decreases. Both

p and q change abruptly at Tjj. The product pq shows relative temperature

independence, reflecting the relatively small bulk-modulus change. The

large temperature- induced changes in p and q present a surprise. It remains

to see whether similar changes occur in nonmagnetic materials.

5. Conclusions

From this study, we reached six conclusions:

1. At low temperatures, in f.c.c. Fe—19Cr—lONi alloys, elasticconstant anomalies arise principally from a paramagnetic—

antiferromagnetic (Neel) phase transition.

2. Adding interstitial carbon-plus -nitrogen to these alloys

expands the unit cell and decreases the Neel temperature.

3. Ducastelle's model, containing only two energy terms with four

parameters, provides a simple, useful approach to describe physical

property changes at magnetic phase transitions.

175
4. Focusing on the repulsive exponential parameter, p, and the

bandstru ture exponential parameter, q, we find that (i) during

cooling co Tfj , p increases approximately 40% and then abruptly

decreases; (ii) similarly, q decreases approximately 30% and then

abruptly increases; (iii) the pq product remains relatively unchanged

versus temperature, reflecting the small bulk-modulus change; (iv)

except for a scaling factor, the q/p(T) ratio looks lika i/(T).

5. Above Tfl, both p and q behave smoothly, failing to reflect the bulkmodulus softening.

6. This simple approach provides a link with more sophisticated (and

complicated) cohesive -energy models where the energy—elastic -constant

relationships remain less transparent.

Acknowledgments

The studies described here represent part of a large group effort.

At the University of South Carolina: C. Almasan, T. Datta, R.D. Edge,

D. Edwards, E.R. Jones. At Oak Ridge National Laboratory: J.W. Cable.

At NIST: M.W. Austin, S.A. Kim, M. Lei.

References

[1] H.M. Ledbetter and M.W. Austin, Mater. Sci. Eng. 70 (1985) 143.

[2] H.M. Ledbetter, M.W. Austin, and S.A. Kim, Mater. Sci. Eng. 85 (1987)

85.

[3] E.R. Jones, T. Datta, C. Almasan, D. Edwards, and H.M. Ledbetter,

Mater. Sci. Eng. 91 (1987) 181.

[4] C. Almasan, T. Datta, R.D. Edge, E.R. Jones, J.W. Cable, and H.

Ledbetter, J. Mag. Mag. Mater., forthcoming.

[5] F. Ducastelle, J. Physique, 31 (1970) 1055.

[6] M.W. Austin and H.M. Ledbetter, unpublished research, NIST, Boulder,

1989.

[7] S. Chickazumi, Physics of Magnetism (Wiley, New York, 1964) p. 82.

[8] H.M. Ledbetter, N.V. Frederick, and M.W. Austin, J. Appl. Phys., 51

(1980) 305.

[9] R. Furth, Proc. Roy. Soc. , A183 (1944) 87.

[10] C.D. Gelatt, H. Ehrenreich, and R.E. Watson, Phys. Rev. B, 15 (1977)

1613.

176
LOW-TEMPERATURE THERMAL EXPANSION OF Fe-Cr-Ni

ALLOYS WITH VARIOUS C+N CONTENTS

M.W. Austin and H.M. Ledbetter

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Institute of Standards and Technology

Boulder, Colorado 80303 (U.S.A.)

Using x-ray diffraction, we measured between 295 and 5 K the

unit-cell size of several Fe—Cr—Ni alloys containing nominally

19Cr and lONi (mass percent) . The alloys contained various

carbon-plus -nitrogen interstitial levels: 0.69 to 1.47 atomic

percent. Two facts motivated our study. First, below 100 K these

alloys undergo paramagnetic—antiferromagnetic transitions that

affect physical properties, especially elastic constants, which

relate strongly to atomic volume. Second, studies on 7- iron

particles (which share the same crystal structure, atomic volume,

and approximate Neel temperature (Tjj) as our alloys) show a 0.3%

volume increase during cooling through the Neel transition. For

the alloys, within our measurement error (0.05%), we failed to

find a volume change at Tjj. Nor did we detect a departure from

cubic symmetry arising from antiferromagnetic spin alignment. All

the alloys fit a simple two-parameter expression for the

dilatation Al/1.

Solids possess three fundamental cohesive properties: cohesive energy,

bulk modulus (reciprocal compressibility), and volume. Small volume changes

produce large physical -property changes. And a small volume change, say

only 0.1%, induces enormous stress:

a - B(AV/V) - 200 GPa • 0.001 - 200 MPa. (1)

This approximates the typical austenitic-steel yield strength (207 MPa)

Such a steel shows an approximate bulk modulus of B - 200 GPa. Volume

changes at structural (first-order) phase transformations : in the f.c.c - b.c.c. martensitic transformation in Fe—31Ni, volume increases 2.6% [1].

Although thermodynamics prohibits volume changes at second- order phase

transitions [2], the V(T) curve may show a discontinuity, thus a change in

thermal expansivity: - (1/V) (dV/dT)

In studying 7- iron particles, Ehrhart and coworkers [3] used x-ray

diffraction to measure volume between 300 and 6 K, through the paramagnetic - antiferromagnetic ordering temperature near 70 K (Tjg) • Cooling showed a

cubic—tetragonal crystal -symmetry change and a 0.3% volume increase.

Intended for publication in Materials Science & Engineering.

177
Fig. 1. Unit-cell size for Fe—20Cr-10Ni alloys with three carbon contents,

shown as mass percent. Arrows indicate Neel transition

temperature determined by d.c. magnetic susceptome try

.

The present report represents a similar study on bulk Fe-Cr—Ni alloys.

In many ways, these alloys resemble 7-iron: (1) same crystal structure;

(2) similar atomic volumes (unit-cell size 3.59 versus 3.58 A [3]); (3)

similar Ndel temperatures (46—87 versus 67 K [3]); (4) similar lowtemperature electron-spin structures [3, 4]; (5) similar neutron-diffraction

atomic moments, 0.55 versus 0.63 jig [9]; (6) similar elastic constants [5].

Fe—Cr—Ni alloys occur in many low- temperature engineering applications

requiring dimensional stability. Thus, the question — does their volume

increase measurably during cooling through Tjj? — contains both practical

and scientific importance.

The studied alloys were already described in detail: their chemical

compositions, elastic properties, and magnetic properties [6—9].

We determined volume -versus -temperature by x-ray diffraction from bulk

specimens. Specimens consisted of 1-mm by 10-mm discs cut from stock with a diamond saw, annealed at 1100°C for 30 min, and electropolished. Each

specimen was mounted inside an evacuated cryostat attached to a horizontal

20-cm-radius Bragg—Brentano x-ray diffractometer. At each temperature, we

scanned specimens continuously at 0.5° (20) min”1

. We used 0.5-mm receiving

slits and 0.1-mm scatter slits. After subtracting background, we leastsquares fitted both Kal and Ka2 peaks to Lorentzean profiles . To correct

for displacement error, we extrapolated a0 (hkl) versus cos8 cot© to 0 - 90°.

At 300 K, this correction was 0.002 A; at 4 K it was 0.005 A.

178
INTENSITY

(arbitrary

units)

Fig. 2. Profile of (200) diffraction peak at three temperatures: above,

near, near Tjj. No peak splitting appears.

179
Figure 1 shows results for alloys with three C+N contents. Curves in

the figure represent least-squares fits to the function

Aa/a0 - a0 + b/[exp(c/T) - 1], (2)

Here a denotes unit-cell size; T absolute temperature; d0 , b, and c

adjustable parameters. This function relates to one used by Swenson [10] to

describe the temperature dependence of silicon's thermal expansivity.

Figure 2 shows intensity profiles of the (200) diffraction peak at

temperatures above and below Tjj. These fail to show evidence of departure

from cubic symmetry.

We need to recognize the obvious question — why, upon cooling through

Tjj, do the Fe—Cr—Ni alloys fail to show a symmetry loss, a volume increase,

and a detectable change in thermal expansivity? Why do the alloys behave

differently from 7- iron, which they resemble strongly?

The remarkable complexity [11] of magnetism in 3d-electron metals

provides several possible answers. Examples include (1) strong sensitivity

of magnetic -exchange energy to interatomic spacing; (2) competing Fe-Fe,

Fe—Ni, Fe—Cr (and so on) interactions; (3) changes in s—d hybridization; (4)

shifts in itinerant—bound electrons; (5) competition among various possible

magnetic states: antiferromagnetic, ferromagnetic, spinglass. After

considering many possibilities, we plan to suggest an answer.

REFERENCES

1. J.F. Breedis and C.M. Wayman, Trans AIME, 224 (1962) 1128.

2. L.D. Landau and E.M. Lifshitz, Statistical Physics, Pergamon, London,

1959, chap. 14.

3. P. Ehrhart, B. Schonfeld, H.H. Ettwig, and W. Pepperhoff, J. Mag. Mag.

Mater., 22 (1980) 79.

4. Y. Ishikawa, Y. Endoh, and T. Takimoto, J. Phys. Chem. Solids, 31

(1970) 1225.

5. H.M. Ledbetter, Phys. Stat. Solid! (a), 85 (1984) 89.

6. H.M. Ledbetter and M.W. Austin, Mater. Sci. Eng., 70 (1985) 143.

7. H.M. Ledbetter, M.W. Austin, and S.A. Kim, Mater. Sci. Eng., 85 (1987)

85.

8. E.R. Jones, T. Datta, C. Almasan, D. Edwards, and H.M. Ledbetter,

Mater. Sci. Eng., 91 (1987) 181.

9. C. Almasan, T. Datta, R.D. Edge, E.R. Jones, J.W. Cable, and H.

Ledbetter, J. Mag. Mag. Mater., forthcoming.

10. C . A. Swenson, J. Phys. Chem. Ref. Data, 12 (1983) 179.

11. W. Marshall (ed.), Theory of Magnetism in Transition Metals, Academic,

New York, 1967

.

180


WELDING PROGRAM

LEADER: T.A. Siewert

CONTRIBUTORS: C.N. McCowan, D.P. Vigliotti

OBJECTIVES

• Investigation of the metallurgical factors that affect the mechanical

properties of stainless-steel weldments at cryogenic temperatures.

• Contribution to the development of improved filler metals for welding

stainless steels for use at 4 K.

• Evaluation of the mechanical properties of weldments at 4 K.

• Development of methods for detecting defects and evaluation of their

significance in stainless steel weldments.

RESEARCH HIGHLIGHTS

• In a study relating inclusions and fracture toughness, type 316L GMA

welds were produced with shielding gases containing a range of oxygen

contents. The strength at 4 K remained constant while the 35%

variation in the fracture toughness was explained by the variation in

the inclusion density. Sections through samples loaded to 85% of

Kic(J) revealed substantial crack growth (0.1 to 0.25 mm).

• A fracture toughness of a type 310 stainless steel weld (25Cr-20Ni) was

evaluated at 4 K. The type 310 weld had substantially poorer toughness

than the other high Ni welds but was similar to that for type 310 base

metal. The toughness was not affected by the inclusion content.

Microstructural evaluation revealed the lower than expected toughness

was due to fracture propagation through an interdendritic phase.

• Substantial effort is being applied to understanding the

micromechanisms of fracture initiation and propagation in stainless

steel welds at 4 K. Specimens are loaded to various levels above and

below Kj c (J) and then sectioned for microstructural analysis. Various

stainless steel alloys behave similarly with void growth being the

minor component of the fracture process and microcrack linkage being

the major component. This is in opposition to most textbooks on

fracture which indicate void growth is the dominant fracture mechanism

(from room temperature studies)

183
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS

Weld Alloy

[designation]

Welding

Process

Tensile Charpy

Impact

Fracture

Toughness

Fatigue Crack

Growth Rate

* Report Volume

(page)

AUSTENITIC STAINLESS STEELS

Fe-13Cr-13Ni-20Mn-0 .2C SMAW XI XI XI (171)

Fe-13Cr-20Ni-9Mn-2Nb SMAW 1,11 I, II. I (159,169)

VIII II (299)

VIII (29)

Fe-15Cr-12Ni-10Mn SMAW IV IV, VIII IV (453)

VIII (29)

Fe-16Cr-15Ni-4Mo SMAW IV IV IV (275)

Fe-16Cr-16Ni-2Mo FCAW II III III, VI III (155)

VI (199)

Fe-16Cr-16Ni-4Mo FCAW IV IV, VIII IV rv (275) j VIII (29)

Fe-16Cr-35Ni-4Mn SMAW I I, VIII II I (159)

[330 modified] II

( (299)

VIII (29)

Fe-17Cr-9Ni-Mn-N SMAW IX IX IX (217)

Fe-17Cr-16Ni-4Mo SMAW IV IV IV (275)

Fe-18Cr-llNi-2Mo-N EBW X X X X (299)

LBW X X X

Fe-18Cr-16Ni-6 . 5Mn-2Mo SMAW IX IX IX (257)

Fe-18Cr-16Ni-9Mn-N GMAW III, IV IV III, IV III, VI III (155)

IV (275,453)

V (199)

Fe-18Cr-18Ni-2Mo-5Mn-N SMAW IV IV IV (453)

Fe-18Cr-20Ni-5Mn-N GMAW IX, XI IX, XI IX (247)

XI (247)

Fe-18Cr-20Ni-6Mn-0.3Nb GMAW IV IV IV IV (275)

Fe-18Cr-20Ni-6Mn-0.2Ti GMAW IV IV IV IV (275)

Fe-19Cr-12Ni-2Mo SMAW II II, VIII II II (315)

[316] VIII (29)

Fe-19Cr-12Ni-2Mo EBW IV IV IV, VI IV (275,415)

[316L] VI (199)

GMAW III, IV III, IV, IV III (155,195)

VIII IV (303)

VIII (29)

MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-884 (I);

NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-1690 (VI);

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX); NBSIR 87-3067 (X); NBSIR 88-3082 (XI)
! MARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS, continued

)

1 Ld Alloy

•signation]

Welding

Process

Tensile Charpy

Impact

Fracture

Toughness

Fatigue Crack

Growth Rate

Report Volume

(page)

-19Cr-12Ni-2Mo GTAW III, IV, III, IV, IV III (155,195)

16L] , continued XI VIII XI IV (303)

VIII (29)

XI (215)

SAW III, IV, III, IV, III, IV III, VI III (155,195

V V IV (275,415)

V (233)

VI (199)

SMAW II, III, II, VIII II, III, III, IV, VI II (299,315)

IV IV III (155,167)

VI (199)

VIII (29)

:-19Cr-12Ni-2Mo-N SMAW II II, VIII II II (299)

I16LN] VIII (29)

:-19Cr-16Ni-7Mn GMAW IV IV IV (453)

*-20Cr-10Ni SMAW I I, VIII I (169)

108] VIII (29)

s-20Cr-10Ni FCAW III III III, VI III (155)

J08L] VI (199)

GMAW III III III, VI III (155)

VI (199)

GTAW XI XI XI (215)

SMAW II, III II, VIII II, III III, VI II (299)

III (155)

VI (199)

VIII (29)

*-20Cr-10Ni-N FCAW IV IV IV, VI IV (275,415)

108LN] V (199)

i-20Cr-10Ni-9Mn-N SMAW I I I (159,169)

»-20Cr-16Ni-7Mo-2W SAW IV IV IV (275)

*-20Cr-34Ni-2Mo-Nb SMAW VII VII VII (253)

>-21Cr-6Ni-9Mn-N GTAW XI XI XI (215)

SMAW I. II I, II I (159,169)

II (299)

1ATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICTIONS AT LOW TEMPERATURES: NBSIR 78-884 (I);

*BSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-1690 (VI);

tfBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX); NBSIR 87-3067 (X); NBSIR 88-3082 (XI)

185
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS, continued

'

Weld Alloy Welding Tensile Charpy Fracture Fatigue Crack Report Volume

[designation] Process Impact Toughness Growth Rate (page)

AUSTENITIC STEELS

Fe-1 5Mn-8Ni-lMo-0 . 7C SAW VIII VIII VIII (15)

Fe-5Cr-25Mn-lNi V V.VII V.VI V (233)

VI (199)

VII (245)

Fe-6 . 5Cr-25Mn-3Ni-lMo SMAW V V V,VI V (233,245)

VI (199)

ALUMINUM ALLOYS

Al-5Mg GMAW III IV rv III (155,217)

[5183] IV (323)

Al-5Mg-0.12Ti GMAW III IV IV III (155, 217)

[5556] IV (323)

Copper Alloys

Cu-0 .08Ag XI XI (201)

[C10700]

Cu-0.ABe-l.8Ni XI XI (201)

[C17510]

*MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICTIONS AT LOW TEMPERATURES: NBSIR 78--884 (i);

NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82 -1667 (V); NBSIR 83-1690 (VI);

NBSIR 8A-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX); NBSIR 87-3067 (X); NBSIR 88-3082 (XI)

186
THE FRACTURE TOUGHNESS OF

25Cr-22Ni-4Mn-2Mo STAINLESS STEEL

WELDS AT 4 K*

C.N. McCowan and T.A. Siewert

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado

A 25Cr-22Ni-4Mn-2Mo stainless steel gas metal arc welding

electrode composition was used to make welds having several

different inclusion contents for fracture toughness (Kj c (J))

testing at 4 K. The inclusion contents of the welds were changed

by increasing the oxygen content of the argon- oxygen shielding gas

mixture. The yield and ultimate strengths of the welds at 4 K

were nearly constant, with values near 875 and 1100 MPa respectively. The fracture toughness of the welds varied from 130 to

165 MPa^m1 / 2

, but this variation in toughness could not be related

to the inclusion content. The metallography and fractography

results show that, although the welds fractured in a ductile mode,

failure occurred predominately at intercellular regions of the

austenitic solidification structure. This intercellular, rather

than a transdendritic , fracture mode is thought to explain the

absence of a strong correlation with inclusion content seen with

other stainless steel welds.

INTRODUCTION

The cryogenic strength and toughness of stainless steel weld compositions that have a primary ferritic (FA) mode of solidification, such as AWS

type 308 and 316 stainless steels, and experimental compositions that have a

fully austenitic (A) mode of solidification, have been evaluated more

extensively than the stainless steel welds which have a primary austenitic

(AF) mode of solidification. 1-5 For welds solidifying in the AF mode,

ferrite forms during the final stages of solidification when the liquid

remaining between the growing austenitic cells becomes enriched to a

eutectic composition. 6-11 This solidification sequence results in a

different composition and ferrite morphology than that occurring in the FA

solidification mode, and the nature of the ferrite-austenite boundary is

changed. These changes could have a substantial effect on the performance

of a weld. This study evaluates the toughness of a modified type 310

stainless steel weld metal that solidifies in an AF mode to determine if

this composition is suitable for cryogenic use. In addition, the effect of

inclusions on the cryogenic toughness of this weld composition are evaluated.

Contribution of the National Institute of Standards and Technology.

187
MATERIALS AMD PROCEDURES

The composition of the 1.2 non diameter stainless steel gas metal arc

(GMA) electrode used to make the four welds in this study is given in

Table 1. The welding was performed at approximately 250A and 22V. The

travel speed was 2.8 mms’ 1 with a contact tip-to-work distance of 19 mm.

An average of 12 weld beads were required to complete the single -vee groove

weldment shown in Figure 1. The base material used was 25-mm- thick type

304L plate. Interpass temperatures were held at about 150 C.

To attain a variation in the inclusion content of the welds, the oxygen

content of the argon- oxygen shielding gas mixture used to make the welds was

changed. To provide baseline data, two welds (ID and IS), were produced

with 100% argon shielding. Weld IS was produced with a standard 19-mmdiameter gas cup but additional shielding was provided for the ID weld by

using a second, larger diameter gas cup which also had a flow of argon. 12

This double shield was used to further reduce the oxygen content in the

atmosphere surrounding the arc. The 1R0 and 1H0 welds were made using

shielding gases having 2 and 10 vol% oxygen respectively and a conventional

19 mm diameter gas cup.

Table 1. The Chemical Composition of the Welding Electrode.

c Si Mn P S Cr Ni Mo Ti Cu Nb Ta

0.02 0.18 4.4 0.010 0.009 24.9 22.3 2.1 0.04 0.02 0.01 0.01

Figure 1. The welding configuration used for the study.

188
A compact tension and a tensile specimen were machined from each of the

welds for testing at 4 K. The locations from which these specimens were

removed from the welds are also shown in Figure 1. The compact tension

specimens were fatigue precracked in liquid N (76 K) and then side grooved

to a depth of 10% of the thickness per side. The fracture toughness testing

was conducted at 4 K (specimens immersed in liquid He) using the single

specimen compliance test method. 8 For the 4 K tensile testing, the specimens were strained at a rate of 2 x 10" 4

. Details of the testing apparatus

are available elsewhere. 13 " 14

RESULTS

A. Microscopy

The location of the ferrite at intercellular regions of the austenite

matrix (Figure 2) identifies the solidification mode of all four welds as

AF. In Figure 3, the intercellular ferrite is shown at a higher magnification: Commonly the ferrite is located at triple points in the cell structure and appears to have nucleated at MnSiO

s -type inclusions. It is evident

that both the ferrite and the inclusions in these welds interact with the

migrating grain boundaries during cooling.

Figure 2. The welds had a cellular austenitic matrix with ferrite present

at the intercellular regions.

189
I

Figure 3. The ferrite often appeared to have nucleated at MnSi03 type

inclusions

.

The types of inclusions observed in the microstructure and the inclusion density (inclusions per unit area) varied with the oxygen content of

the welds. At low oxygen contents, spherical, MnSiO

s type inclusions and

Ti-rich inclusions were most common (Figure 4). As the oxygen content of

the welds increased (1R0 and 1H0) , the number of MnSi03 and Ti-rich inclusions per unit area increased, and large, Cr-rich inclusions were also

frequently observed. The large number of Cr- and Ti-rich inclusions

observed was unexpected and precluded the collection of inclusion density

counts for comparison to previously collected inclusion data. It is clear

in Figure 5, however, that the welds had very different inclusion contents

and so any effect of inclusion content on the toughness should be apparent.

B. Mechanical Properties

The 4-K tensile and fracture toughness results for the welds are given

in Table 2. The average yield strength, ultimate strength, and fracture

toughness (Kj c (J)) values of the welds were 878 MPa, 1104 MPa, and

143 MPa-m^ respectively. The yield and tensile strength are probably

underestimated by these averages, however, because the low values reported

for the IS weld are attributed to a solidification crack in the specimen.

Several other notes concerning the missing data in Table 2 are: (1) the 1R0

tensile specimen was accidentally tested at a very high strain rate (7.5 x

10" *) and failed prior to reaching a 0.2% offset yield strength, and (2) no

fracture toughness value is reported for the IS weld because the test was

stopped at the first sign of crack growth so that the specimen could be

sectioned for metallographic examination of the crack front region.

190
Figure 4. The spherical inclusions are MnSi03 type inclusions typically

found in stainless steel welds. The more angular, dark inclusions were rich in Cr and/or Ti.

The strength and toughness values determined for the welds in this

study are representative of the type 300 series stainless steel welds. In

Figure 6, for example, the strength and toughness of the 25Cr-22Ni welds

were similar to those found for a previously tested AISI 310 weld (25Cr21Ni-2Mn-0 . 10C) and are near the middle of the strength- toughness zone

determined for type 308 and 316 stainless steel welds. 1 » 15 » 16 Unexpectedly,

no trend of increasing toughness with decreasing inclusion contents in the

welds was found. The weld having the highest inclusion content (1H0) had a

toughness similar to the weld with the lowest inclusion content (ID).

Table 2. The 4 K tensile and fracture toughness results.

Weld

Yield Strength

(MPa)

Ultimate Strength

(MPa)

Elongation

(%)

Reduction in

Area (%)

Kic(J)

(MPa/m)

ID 907 1178 15.6 9.6 128.6

IS 836 907 2.8 8.0 —

1R0 — — — — 164.8

1H0 890 1227 24.1 14.5 135.2

191
(b)

(c)

Figure 5. The inclusion distributions for the (a) ID weld, (b) the 1R0

weld, and (c) the 1H0 weld are shown.

192
Figure 6. The strength and toughness of the ID and lHo welds are compared

here with previous data collected for type 300 series stainless

steel welds and base materials.

Based on the difference in the inclusion contents of the ID and 1H0

welds (Figure 5), but similar toughness, inclusion content does not appear to

control the crack initiation in these weld specimens. Since previous studies

have shown a similar change in inclusion content to cause a 35% change in

toughness at 4 K, 12 we evaluated fracture surfaces and metallographic

sections from the weld specimens to determine the cause for this anomalous

behavior

C . Failure Analysis

The fracture surfaces of both tensile and compact tension specimens

indicated that failures occurred in a ductile manner, but the fracture path

was determined by the microstructure. The fracture surface of the ID tensile

specimen, for example, is shown in Figure 7 to be dominated by ductile

dimples that form a pattern of peaks and valleys. Areas having these ribbed

patterns were observed to some extent on each of the tensile fracture

surfaces and were recognized as being similar to the solidification cell

dimensions shown in Figure 3. This peak-valley pattern could be generated if

fracture occurred preferentially through the intercellular regions of the

solidification structure.

193
Figure 7. The fracture surface of the ID tensile specimen showed a ductile

dimple network that formed a series of peaks and valley.

Further evidence of intercellular fracture was found during examination

of the compact tensile specimens. A transverse section through the initial

crack growth region in the IS compact tension specimen, Figure 8, produced

fracture surface profiles that closely resemble the austenitic cell structure. When examining the crack path, numerous peaks and valleys can be found

that terminate at intercellular positions in the matrix and the end of the

crack tip is located at a cell boundary. In addition, voids observed in the

matrix are located exclusively at intercellular regions, most are at

triple points of the austenitic cell matrix. Because the fracture toughness

test was stopped at crack initiation for the examination of this specimen, a

Kj c (J) value could not be determined. The applied load at crack initiation,

however, was within 10% of those estimated for crack growth in the ID and

1H0 test specimens, so the fracture toughness would be expected to be comparable.

The observations of void formation at triple points in the austenitic

cell structure suggested that the ferritic phase, often present at these

sites, was playing a key role in fracture initiation. Further examination

of the fracture surfaces of the compact tension specimens revealed that

ferrite was present at many of the void sites, but whether the voids

initiated at ferrite -austenite interfaces, in the ferrite, or in the

austenite was not clear. Metallographic cross sections of deformed tensile

specimens (Figure 9), however, did indicate that: (1) the ferrite was

fracturing in a brittle manner as plastic flow occurred in the austenitic

matrix, (2) the ferrite -austenitic interface was generally intact, (3)

cracks sometimes linked neighboring ferrite particles, and (4) several

cracks, apparently nucleated at inclusions in the cell cores, were observed

to change growth directions and follow along the cell boundary upon reaching

the intercellular regions. These observations suggest that the probable

cause of void initiation at ferrite particles is a result of the brittle

fracture of the ferrite.

194
Figure 8. The crack path for the IS compact tension specimen is shown to

follow the intercellular regions of the microstructure.

Figure 9. In the IS tensile specimen, deformation of the austenitie matrix

often resulted in the fracture of ferrite particles and sometimes

cracks at inclusion sites were observed. The textured appearance

of the matrix is due to plastic flow.

195
These observations, except that of the apparently strong ferriteaustenite interface, are different from those made in deformed ferrite plus

austenite stainless steel welds that have solidified in an FA solidification

mode. In a previous study of type 316L weld tensile specimens that solidified in an FA mode, ferrite was not found to fracture nor were cracks at inclusions or between ferrite particles observed after 50% elongation at 4 K. 12

For the welds in this study, the AF solidification mode seems to have

produced ferrite and intercellular compositions that limited the toughness.

In a 22Cr-14Ni weld, for example, Brooks estimated the intercellular

compositions to be enriched in Ni and Cr by approximately 2 wt.%: The

ferrite composition was approximately 34Cr-8Ni. 9 At nominal alloy contents

farther from the liquidus trough of the Fe-Cr-Ni ternary diagram than

Brook's alloys, such as the 25Cr-22Ni welds in this study, even more

extensive differences between the core and intercellular compositions could

occur. In addition, the welds evaluated in this study would be expected to

have a Mo -enriched ferrite phase and intercellular regions that are enriched

in Mo and Mn. Unfortunately, the detailed microanalysis required to

quantitatively determine the level of microsegregation in these welds was

beyond the scope of this study. The microsegregation, however, did not

produce cell core chemistries that transformed to martensite during deformation at 4 K, so the corduroy type fracture mechanism reported by Lippold is

not applicable to these welds. 17

CONCLUSIONS

1. The fracture toughness values of the 25Cr-22Ni-4Mn-2Mo at 4 K were

typical of those that have been determined for type 308, 316, and 310

stainless steel welds of similar strength and average inclusion

contents

.

2. The failure of the tensile and compact tension specimens occurred

predominately at intercellular regions in the weld microstructure.

3. The failures are suspected to initiate at local brittle fractures in

the ferrite, which then link along intercellular paths to form a crack.

ACKNOWLEDGMENTS

The electrode was furnished by Sandvik Steel Company.

REFERENCES

1. Tobler, R.L.

,

Siewert, T.A.

,

and McHenry, H.I., 1986, StrengthToughness Relationship for Austenitic Stainless Steel Welds at 4 K,

Cryogenics, July, pp. 3922-395.

2. McCowan, C.N., Siewert, T.A.

,

and Tobler, R.L. 1986, Tensile and

Fracture Properties of an Fe- 18Cr-20Ni-5Mn-0 . 16N Fully Austenitic Weld

Metal at 4 K, Journal of Engineering Materials and Technology, October,

pp. 340-343.

196
3. Elmer, J.W., McHenry, H.I., and Whipple, T.A.

,

1981, Strength and

Toughness of Fully Austenitic Stainless Steel Filler Metals at 4 K,

in: Materials Studies for Magnetic Fusion Energy Applications at Low

Temperature - IV, NBSIR 81-1645, National Bureau of Standards, Boulder,

Colorado, pp. 289-302.

4. Tobler, R.L.

,

Trevisan, R.E. and Reed, R.P., 1985, Tensile and Fracture

Properties of an Fe-14Mn-8Ni-lMo-0.7 Fully Austenitic Weld Metal at

4 K, Cryogenics, Volume 25, pp. 447-451.

5. Yamagami, N. , Kahsaka, Y. , and Ouchi, L. , 1987, Mechanical Properties

of Welded Joints In Mn-Cr and Ni-Cr Stainless Steels at 4 K, in:

Advances in Cryogenic Engineering - Materials, Clark, A.F. and Reed,

R.P., Ed., Plenum Press, Volume 34, pp. 359-366.

6. David, S.A, Goodwin, G.M., and Braski, D.N., 1979, Solidification

Behavior of Austenitic Stainless Steel Filler Metals, Welding Journal,

Nov., pp. 330s-336s.

7. Ritter, A.M. Henry, M.F., Savage, W.F., 1984, High Temperature Phase

Chemistries and Solidification Mode Prediction in Nitrogen- Strengthened

Austenitic Stainless Steels, Metallurgical Transactions A, July,

pp. 1339-1351.

8. Cieslak, M.J., Ritter, A.M, and Savage, W.F. , 1982, Solidification

Cracking and Analytical Electron Microscopy of Austenitic Stainless

Steel Weld Metals, Welding Journal, January, pp Is -8s.

9. Brooks, J.A., Williams, J.C., and Thompson, A.W. , 1983, STEM Analysis

of Primary Austenite Solidified Stainless Steel Welds, Metallurgical

Transactions A, January, pp 23-31.

10. Takalo, T. , Suutala, N. , and Moisio, T. , 1979, Austenitic Solidification Mode in Austenitic Stainless Steel Welds, Metallurgical Transactions A, August, pp 1173-1181.

11. Kokawa, H. , Kuwana, T. , and Yamamoto, A., 1988, Crystallographic

Characteristics of Delta- Ferrite Transformations in Type 304L Austenitic Stainless Steel Weld Metal at Elevated Temperatures, IIW Doc

IX- 1529-88

12. McCowan, C.N. and Siewert, T.A. , 1989, The effect of Inclusions on the

Fracture Toughness of 316L Stainless Steel Welds at 4 K, submitted to:

Advances in Cryogenic Engineering - Materials.

13. Tobler, R.L. Read, D.T.

,

and Reed, R.P., 1981, Strength/Toughness

Relationship for Interstitially Strengthened AISI 304 Stainless

Steels at 4 K, in Fracture Mechanics: Thirteenth Conference, Richard

Roberts, Ed. ASTM STP-743, American Society for Testing and Materials,

Philadelphia, pp. 250-268.

197
14. Read, D.T. and Tobler, R.L. , 1982, Mechanical Property Measurements at

Low Temperatures, in Advances in Cryogenic Engineering - Materials,

Vol. 28, Plenum Press, New York, pp. 17-28.

15. Tobler, R.L. , Shoji, T. , Takahashi, and Ohnishi, K. , 1986, Fracture,

Acoustic Emission and Adiabatic Heating of Austenitic Stainless Steels

at Liquid Helium Temperature, Yamaguchi, K, Aoki, K. , and Kishi, T.

,

Ed., in Progress in Acoustic Emission III, The Japanese Society of NDI,

pp. 453-461.

16. Reed, R.P.

,

Simon, N.J., Purtscher, P.T., and Tobler, R.L.

,

1986, Alloy

316 LN for Low Temperature Structures: A Summary of Tensile and

Fracture Data, in Proceedings of the Eleventh International Cryogenic

Engineering Conference, Kipping, G. and Kipping, I., Ed., Butterworth

and Co. Surrey UK, pp. 786-793.

17. Lippold, J.C., Juhas, M.C., and Dalder, E.N.C., 1985, The Relationship

between Microstructure and Fracture Behavior of Fully Austenitic Type

316L Weldments at 4.2 K, Metallurgical Transactions A, October, pp.

1835-1848.

198
THE EFFECT OF INCLUSIONS ON THE FRACTURE TOUGHNESS

OF 316L STAINLESS STEEL WELDS AT 4 K

C.N. McCowan and T.A. Siewert

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado

The 4-K fracture toughness (Kj c (J)) of a type 316L stainless

steel weld composition increased significantly when the inclusion

contents of the welds were decreased. For the three welds tested

(Ni - 13 wt.%, 4 K yield strength » 740 MPa), the fracture

toughness increased from 132 to 150 to 179 MPa-mh as the inclusion

contents decreased from 55,200 to 37,700 to 19,300 inclusions per

mm2

. Expressing the effect of inclusions on the toughness in

terms of the inclusion spacing (inverse square root of the

inclusion density) , the fracture toughness of the 316L welds

increased 18 MPa*mH per /im increase in the average inclusion

spacing. Previous results from a 18Cr-20Ni stainless steel weld

study are expanded on in this study and show a trend similar to

the trend found for the 316L welds. For the 18Cr-20Ni welds (4-K

yield strength * 1000 MPa) , the fracture toughness increased from

180 to 287 MPa*mh as the inclusion contents decreased over the

range of 25,000 to 8,300 inclusions per mm2

. Because the effects

of yield strength and Ni content on the toughness of the two weld

compositions were constant, the increases in toughness (35% for

the 316L, 60% for the 18Cr-20Ni) are attributed solely to the

reductions in weld metal inclusion contents. For the 316L welds,

fractographic studies were conducted to characterize differences

in the initiation of fracture as a function of the inclusion

content. Measurements on fracture specimens loaded to only 85% of

the estimated Kj c values revealed stable crack growth lengths of

0.1 to 0.25 mm crack growth occurs as voids which formed at

inclusions link by microcracking, rather than coalescence.

INTRODUCTION

Three variables that affect the fracture toughness of stainless steel

welds and base materials at cryogenic temperatures are yield strength, Ni

content, and inclusion content. 1 " 3 The toughness (Kj c (J)) increases with

increasing Ni content, and decreases with increasing strength and inclusion

content. Therefore, it is not surprising that welds, which generally have

higher inclusion contents than base materials, are typically found to have a

199
lower toughness than base materials of comparable strength and Ni levels.

But the importance of inclusions in determining the toughness is not well

defined for stainless steel welds and the quantitative dependence of the

previously observed trends on the strength level, nickel content, and microstructure of an alloy has not been determined. This information is needed to

help determine what increases in the toughness of welds at cryogenic can be

expected by improved inclusion control.

The effect of inclusion content on the fracture toughness at 4 K has

been studied in an 18Cr-20Ni weld composition. 1 In that study the fracture

toughness (Kj c (J)) increased by 30 to 40% when inclusion contents were

decreased from approximately 11,000 to 9,000 inclusions per mm2

. This study

adds to the work previously done on the 18Cr-20Ni GMA welds and evaluates the

effect of inclusion content on the fracture toughness of type 316L stainless

steel gas metal arc (GMA) welds.

MATERIALS AND PROCEDURES

The chemical compositions supplied by the manufacturers of the two GMA

electrodes used in this study are given in Table 1. The oxygen contents of

the welds (Table 2) made for the study were measured with a fusion analysis

technique using a thermal conductivity detector.

To produce welds having different inclusion contents, the oxygen content

of the argon-oxygen shielding gas and the method of shielding used to make

the welds were varied. Welds having inclusion contents most typical of

commercially produced stainless steel GMA welds were made using an argon- 2%

oxygen (by volume) shielding gas and a conventional 16 mm (5/8 inch) gas

nozzle. To attain higher inclusion content welds, an argon based shielding

gas having approximately 8% oxygen was used. For low weld inclusion contents, a dual -gas -shielding configuration (Figure 1) was implemented using

100% argon for both the inner and outer shields. The shielding gas compositions were controlled during welding by using mass flow meters to mix

commercial grade oxygen and argon. The total flow rate of the shielding gas

was 14.2 L per min (30 cfh)

.

Table 1. The GMA electrode compositions in wt.%.

C Mn Si P S Cr Ni Mo N Cu

316L 0.02 1.73 0.35 0.008 0.009 19.2 13.1 2.15 NA 0.04

18Cr-20N 0.03 5.39 0.26 0.006 0.007 18.1 20.4 0.11 0.16 0.02

200
able 2. The inclusion and 0 2 and contents of the welds. The D, Ro, and Ho

suffixes stand for double Argon, Argon - 2% Oxygen, and Argon - 8% Oxygen shield respectively.

Weld o 2 Inclusion

Density

Inclusion

Spacing

316L-D 0 . 0004 19,300 7.0

316L-Ro 0.048 37,700 5.0

316L-Ho 0.072 55,200 4.3

11 0.005 9,800 11.0

12 0.004 8,300 10.1

10 0.037 11,300 9.4

13 0.056 25,000 6.3

igure 1. The dual-shield configuration used for this study to provide better

shielding of the welding arc.

201
The specific welding parameters used for the welds are included in Table 3.

In general, we tried to maintain the same heat input, contact tip -to -work

distance and interpass temperature for all welds made with the same electrode.

The exception was the 18Cr-20Ni number 11 weld that was made as part of a

previous study. 1 For this weld, the welding parameters were altered to change

the metal transfer mode from a spray to droplet type.

From each weld, a uniaxial tensile and two compact tension specimens were

machined for testing at 4 K. The weld configuration and specimen locations are

shown in Figure 2. The compact tension specimens were precracked in liquid

nitrogen and then side grooved on each side to a depth of 10% of the specimen

width prior to testing. The fracture toughness tests were conducted with the

specimens submerged in liquid helium (4 K) using the single specimen compliance

method. The uniaxial tensile tests were also conducted at 4 K: tensile specimens having diameters of 6.25 mm were strained at a rate of 2 X 10" 4

, the percent

elongation was measured over a 25.4 mm gauge length, and the yield strength was

calculated using the 0.2% offset method. The details of the 4-K testing procedures are given elsewhere

.

4 » 6

Table 3. Welding and Testing Parameters used for the welds.

Volts Amperage Electrode

Extension (mm)

Electrode

Diameter (mm)

316L-D-RO-HO 24 210 19 0.9

10*, 12*, 13 25 280 18 1.2

11* 20 175 12 1.2

Welded as part of a previous study. 1

Figure 2

.

202
To estimate the inclusion densities, cross sections of the velds

perpendicular to the welding direction were cut and polished for scanning

electron microscope (SEM) evaluations. For each weld, 200 to 300 inclusions

were counted at 2000x and then divided by the area scanned to determine the

inclusion density. The average inclusion spacing was calculated as the

inverse square root of the inclusion density. 6

To evaluate damage to the weld microstructures near the onset of stable

crack growth, the second fracture toughness test for each of the 316L welds

was interrupted at 85% of the estimated Kjc

(J) value so that the specimens

could be removed and examined. Samples were prepared for light and scanning

electron microscope evaluations as follows: (1) samples containing all of the

fatigue crack and a reasonable volume of the crack front region were removed

using an abrasive saw, (2) these samples were mounted in epoxy under vacuum,

(3) by rough grinding, a plane perpendicular to the fatigue crack plane and

parallel to the direction of crack growth was obtained at a depth of approximately 5 /im from the original compact tension surface, and (4) following

final polishing and observations on this plane, samples were reground and

repolished numerous times to provide a series of sample planes from different

depths within the specimens.

RESULTS AND DISCUSSION

Microstructure

The 316L and the 18Cr-20Ni welds had different solidification modes. The

18Cr-20Ni, because of the higher Ni and N content, solidified in a fully

austenitic mode. The 316L welds solidified in a primary ferritic mode. The

residual ferrite contents measured for the 316L welds using magnetic response

techniques in accordance with ANSI/AWS Standard A4.2-74 ranged from 5 to 7 FN.

The variations in the inclusion contents observed for the 316L welds are

shown in Figure 3 and estimated inclusion densities for all the welds are

given in Table 2. Clearly a significant range in inclusion contents was

achieved for both weld compositions: the percent increase in inclusions for

the 316L and the 18Cr-20Ni welds over their respective ranges were 185 and

200%. The majority of the inclusions observed in the welds were spherical,

MnSiOs type inclusions. Occasionally, more angular inclusion shapes were

observed that were determined during energy dipersive x-ray analysis to be

rich in Cr, Ti, and sometimes Ai.

Mechanical Testing

The yield strength of the welds varied by less than 4% for a given

electrode composition (Table 4) so the effect of strength and Ni on the

toughness are essentially constant for the 316L and 18Cr-20Ni welds. The

average yield strength was 742 MPa for the 316L welds and 1005 MPa for the

18Cr-20Ni welds. There was no correlation between the inclusion density and

the yield strength of the welds.

203
Figure 3. The inclusion content range for the 316L welds: (a) 316L-D and

(b) 316-HO. These SEM micrographs were taken on polish planes

perpendicular to the welding direction.

204
Table 4. The mechanical properties of the weld at 4 K.

°y 0TTs %E %R Kic(J)

316L D 736 1343 47.9 46.6 179

316L Ro 747 1153 22.3 23.7 150

316L Ho 743 1122 10.2 13.1 132

10 991 1274 23.0 287

11 1027 1347 17.0 15 263

12 1015 1300 26.0 37.0 203

13 988 1209 7.7 16.6 180

The ultimate strength, elongation, and reduction in area for the tensile

specimens are included in Table 4. Although there appears to be a trend of

decreasing tensile ductility (as measured by elongation and reduction in

area) with increasing inclusion content for the 316L welds, no such trend is

evident for the 18Cr-20Ni welds. In part, this apparent difference in the

tensile ductility of the welds is due to scatter in the data. Weld specimens, because of their coarse dendritic crystal structures, develop heavy

surface textures during straining at 4 K that can result in stress concentrations which initiate failure. Therefore, variation in the tensile ductility

data may not be due to the inclusion content. Alternately, microflaws in the

weld specimens sometimes initiate premature failures.

The fracture toughness (Ki c (J)) increased as the inclusion content of

the welds decreased (Table 4). For the type 316L stainless welds, the

fracture toughness increased about 35% when the inclusion density was lowered

from 55,200 to 19,300 inclusions per mm2

. This change in toughness with

inclusion content is shown in Figure 4 to almost span the width of the

strength- toughness trend band previously developed for type 316L and 308L

stainless steel welds at 4 K. 7 The trend band was constructed using available stainless steel weld data and the width of the band were attributed to

variations in weld inclusion contents between the different welding processes, which strongly supports this observation. The type 316L stainless

steels welds tested in this study have an inclusion range that corresponds

reasonably well to the range produced by common production welding processes

and strongly supports this observation.

205
Figure 4. The 316L (stars and the 18Cr-20Ni (circles) weld data show

increases in the toughness with decreasing inclusion contents.

The changes in toughness are of a similar magnitude to the trend

bandwidth determined for type 308 and 316 stainless steel welds

made by a variety of commercial welding processes.

One way to express the effect of inclusions on toughness is in terms of

the average inclusion spacing (inverse square root of the area density).

This approach has been used to model fracture initiation. It provided a

characteristic distance (inclusion spacing) over which a critical stress or

strain must extend for the onset of stable crack growth. 8 The results of

this study (Figure 5), as expected, show that the toughness increases

significantly with inclusion spacing for both the 316L and the 18Cr-20Ni

welds. Although the trends indicated by these data are not expected to be

linear over a large range, the 316L weld data shows the fracture toughness at

4 K to increase by approximately 18 MPa-mh per pm increase in the inclusion

spacing. A trend of similar magnitude is seen for the 18Cr-20Ni data. No

differences in the trends for the two types of welds are due to differences

in the inclusion spacing ranges, solidification mode, strength, and Ni

content of the weld were resolved. As more data is collected for welds,

these effects should become apparent.

In Figure 6 , the effects of strength and Ni content on the toughness are

normalized so that the effects of inclusion spacing indicated for the two weld

compositions can be more fairly compared with each other and with type 316L

stainless steel base materials at 4 K.® Considering that the base metal trend

was developed at inclusion spacings greater than 50 /*m, the agreement of the

extrapolated trend with the weld data is surprisingly good. All of the weld

data is within the scatter band associated with the base metal trend.

206
I

a

Figure 5. The increases in the toughness for the 316L (stars) and the 18Cr20Ni (circles) welds are shown here in terms of increasing average

inclusion spacing (inverse square root of the inclusion density).

5 xio3

Figure 6. The Quality Index (Kj c -ay), normalized to a 12 wt.% nickel content,

versus the inclusion spacings of the welds.

207
Although extrapolation of these weld data trends would indicate superior

toughness for the welds compared to base materials at inclusion spacing

greater than 12 to 15 /im, inclusion size and morphology are not considered

here, so such speculations are not appropriate. Perhaps the most interesting

comparison to be made in Figure 6 is between the two types of solidification

structures present for the welds. Unexpectedly, the fully austenitic structure of the 18Cr-20Ni welds does not appear to be inherently tougher than the

2 -phase ferrite plus austenite structure of the 316L welds.

Failure Analysis

All test specimens failed in a ductile manner with fracture surfaces

exhibiting typical ductile dimple morphologies (Figure 7) . Evaluations of the

316L specimens that were only loaded to 85% of their estimated Kj c (J) values

showed more crack growth than expected (Figure 8) . To examine the initiation

of crack growth, we estimate that the fracture tests should have been interrupted at 70% of the respective Kj c (J) values. Still, there were several

characteristic differences observed in the appearances of the 316L crack

paths. The lower inclusion content specimen (316L-D) typically had an area of

fracture out in front of the fatigue crack that appears to have formed first

and then linked by microcracking to the fatigue crack tip. Sharp crack tips

were commonly observed at the leading edges of these ductile fracture zones.

For the 316L-RO and -HO specimens, a more continuous or shorter crack voidlinking distance was observed ahead of the fatigue crack.

Figure 7. The fractures of the compact tension and uniaxial tensile specimens

were dominated by ductile dimples (316L-HO compact tension specimen) .

208
120 pan 1

Figure 8. Stable crack growth was observed in each of the 316L fracture

specimens at 85% of their estimated Kjc values. The SEM micrographs shown here are from the 316L-D (A) , the 316L-RO (B) , and the

316L-HO (C) compact tension specimens.

The void sizes observed on cross-sections of the tensile and compact

tension samples are small compared with the ductile dimple sizes on the

fracture surfaces (Figure 9) . This implies that much of the dimple growth

process is due to localized flow during fracture: voids do not appear to

grow stably until impingement causes failure on a general basis. Most

commonly, voids were observed to be linked by "microcracks" during the

examination of the 316L crack profiles (Figure 10) . Although the "microcracks" are sometimes closely associated with the ferrite phase (Figure 11),

our observations indicate that this is due to a second order effect of the

ferrite phase interrupting crack growth directions in the softer austenite,

rather than the ferrite -austenite initiating cracks. Voids appear to

nucleate at precipitates or inclusion associated with the ferrite phase. The

ferrite phase, however, does serve as an effective barrier to plastic flow in

the matrix (Figure 12)

,

and this may be the principal cause of the increased

strength and reduced toughness that has been attributed to the presence of

ferrite in austenitic welds.

209
Figure 9. The voids that nucleated near the crack planes in the 316L

fracture specimens typically showed only limited signs of growth

(316L-HO sample, 85% Kjc )

.

CONCLUSIONS

1. Improved gas shielding of the arc region during GMA welding results in

welds of significantly increased toughness at 4 K.

2. Type 316L stainless steel welds already contain appreciable crack growth

at 85% of their estimated Kjc (J) values.

3. Stable crack growth appears to occur by the linking of voids via "microcracking" rather than the coalescence of voids.

210
Figure 10. The crack tip region of the 316L-HO specimen (85% of Kjc) shows

voids to be linked by "microcracking" rather than void coalescence. For Figure 10A, a 20 ml HC1, 20 ml lactic acid, 5 ml HN0 3 was etched.

211
Figure 11. The ferrite phase (medium contrast, skeletal network) was sometimes found to be closely assoicated with the fracture path (316LHO sample, 85% of Kjc )

.

Figure 12. In the 316L-D tensile specimen (48% elongation)

,

ferrite was observed. The etch (20 ml HC1 , 20 ml

ml HN0 3 ) shows ferrite (white) and plastic flow.

no fractured

lactic acid, 5

212
REFERENCES

1. McCowan, C.N. and Siewert, T.A. , Inclusions and Fracture Toughness in

Stainless Steel Welds at 4 K, Advances in Cryogenic Engineering -

Materials, Volume 32, 1986, Plenum Press, 335-342.

2. Morris, J.W., Structural Alloys for High Field Superconducting Magnets,

Advances in Cryogenic Engineering-Materials, Volume 32, 1986, Plenum

Press, 1-22.

3. Reed, R.P., Purtscher, P.T., and Yushchenko, K.A.

,

Nickel and Nitrogen

Alloying Effects on the Strength and Toughness of Austenitic Stainless

Steels at 4 K, Advances in Cryogenic Engineering-Materials, Volume 32,

1986, Plenum Press, 43-50.

4. Tobler, R.L.

,

Read, D.T.

,

and Reed, R.P., Strength/Toughness

Relationship for Interstitially Strengthened AISI 304 Stainless Steels

at 4 K, Temperature, Fracture Mechanics: Thirteenth Conference, Richard

Roberts Ed., ASTM STP 743, America Society for Testing and Materials,

Philadelphia, 1981, 250-268.

5. Standard Test Method for Jj c , A Measure of Fracture Toughness, E813-81,

in: Annual Book of ASTM Standards, Part 10, American Society for

Testing and Materials, Philadelphia, 1982, 822-840.

6. Underwood, E.E., Quantitative Stereology, Addison-Wesley Publ. Co,

Reading, MA, 1970.

7. Tobler, R.L., Siewert, T.A. and McHenry, H.I., Strength -Toughness Relationship for Austenitic Stainless Steel Welds at 4 K, Materials Studies

for Magnetic Fusion Energy Applications at Low Temperatures - IX (NBSIR

86-3050), 1986, 239-246.

8. Ritchie, R.O. and Thompson, A.W.

,

On Macroscopic and Microscopic

Analyses for Crack Initiation and Crack Growth Toughness in Ductile

Alloys, Metallurgical Transactions A (16A) 2, 1985, 233-247.

9. Reed, R.P., Simon, N.J., Purtscher, P.T.

,

and Tobler, R.L.

,

Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures — IX

(NBSIR 86-3050), 1986, 15-26.

213

TEST STANDARDS

215

UNITED STATES -JAPAN DEVELOPMENT OF TEST STANDARDS PROGRAM

PRINCIPAL INVESTIGATOR: R.L. Tobler

CONTRIBUTING ORGANIZATIONS: National Institute of Standards and Technology,

Tohoku University,

Japan Atomic Energy Research Institute,

Kobe Steel,

National Research Institute for Metals

OBJECTIVES

• Establishment of a standard test method for tensile testing of structural

alloys at 4 K

• Establishment of standard test methods for fracture testing at 4 K

• Administration of a round-robin test program to support the standards

development

• Administration of experimental research to evaluate critical aspects of

test standardization at 4 K

• Liaison with the American Society for Testing and Materials (ASTM) and

its task group activities for the purpose of establishing the standards

by voluntary consensus

PROGRAM HIGHLIGHTS

• Cryogenic materials scientists were interviewed at government, industrial, and academic institutions in Japan. Their experience and advice

were helpful in preparing the drafts of the test procedure standards.

• Six drafts of a proposed standard 4-K tensile test were written and

submitted for review at domestic and international workshops. The sixth

revision has been formally submitted to ASTM, and is being balloted.

• Five drafts of a proposed standard for 4-K fracture toughnes testing were

written and reviewed at Japanese domestic workshops . A sixth revision is

being prepared for ASTM to begin due process for standardization.

• Round- robin tensile and fracture tests were conducted on a Fe-24Cr-15Ni

alloy. A summary of results was presented at the 1988 United States

Japan International Workshop. A paper has been submitted for publication

at the International Cryogenic Materials Conference to be held July 24-

28, 1989 in Los Angeles.

217
• An experimental research program was completed by participants at NIST

and four Japanese laboratories. From 70 Jj c fracture toughness tests,

the effects of specimen size, side grooves, test speed, and fatigue

cracking on the test procedure were determined. Twelve papers describing

the results have been published.

• Six workshops were an integral part of the program:

Standardization of Fracture Toughness Testing of Low- Temperature

Structural Materials at Tohoku University, Sendai, Japan in March 1986.

Low -Temperature Structural Materials and Standards at Japan Atomic

Energy Research Institute in Tokai-mura, Japan in August 1986

DOE-JAERI Low Temperature Structural Materials and Standards Workshop

in Reno, Nevada in October 1986

United States -Japan Cooperative Research Program Review at Tohoku

University in January 1987

United States -Japan Cooperative Research Program Review at the Japan

Atomic Energy Research Institute, Tokai-mura in May 1988.

Workshop on 4-K Mechanical Testing at the National Institute of

Standards and Technology, Boulder on April 4, 1989.

• At ASTM technical meetings, the draft standards are being reviewed by the

task groups that will process these standards. The following meetings

helped to advance these standards:

The E24 meeting on fracture testing held in Reno Nevada, April 16-17,

1988.

The E28 meeting on tensile testing held in Baltimore, Maryland May 16-

17, 1988.

The E24 and E28 meetings held in Atlanta, Georgia, November 6-9, 1988.

The E24 and E28 meetings held in Kansas City, Missouri, May 22-25,

1989.

218
U.S. - JAPAN COOPERATIVE PROGRAM FOR THE

DEVELOPMENT OF 4-K TEST STANDARDS

Technical meetings (and follow-up discussions)

R.L. Tobler

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado

INTRODUCTION

With the American Society for Testing and Materials (ASTM) , we are

establishing standards for tensile and fracture toughness testing of alloys

at 4 K. We are also interested in the status of standards for testing

composite materials. In support of these activities, I participated in the

following meetings:

1. The U.S. domestic workshop

Colorado Springs, Colorado, October 18-19, 1988, and

2. ASTM Technical Committee Week,

Atlanta, Georgia, November 7-10, 1988.

Since these meetings, several follow-up discussions have taken place. The

purpose of this memorandum is to bring you up to date on the progress of all

the related standards activities.

THE 1988 U.S. DOMESTIC WORKSHOP

At this meeting I described some technical highlights from the U.S. -

Japan international workshop held at JAERI May 27-28, 1988. I then

presented a preliminary comparison of the second round robin results from

the U.S. laboratories.

Four U.S. laboratories participated in the round robin testing of JN

steel. Discussion focussed on the fracture toughness measurements because

several problems were encountered. The data from two laboratories show

unacceptable scatter in resistance curves and significant bias in Kj c (J)

measurements; the values from one lab are too low (about 100 MPa-m^), and

the values from another are too high (about 300 MPa-m

1*).

All four laboratories used the unloading compliance technique. The

system at NIST has been used since 1980. The other three systems were

developed very recently, and two of them may still be bugged. Someone

suggested that material variability could have contributed to the scatter,

but there is no evidence of such variability in the Japanese results which

are all very uniform.

219
In January 1989, JAERI and NIST agreed to publish the results of this

round robin at the next meeting of ICMC (July 24-28, 1989). Therefore we

have about five more months to resolve the U.S. results. After separate

discussions with representatives of each lab (I. Hwang, MIT; L. Summers,

LLNL; and J.W. Morris, LBL) , I recommend the following actions to explain or

improve the results:

1. Each U.S. participant will review their procedure and results

to be sure no obvious mistakes were made

.

2. Each participant will send fractured specimens to LBL for

microscopic examination to verify that there are no anomalous

fractographic features reflecting material variability.

3. The participants will consider a meeting of appropriate

personnel to discuss programming and compare flow charts of

their data acquisition and reduction methods.

Rer "ding the third point above, NIST can also benefit by upgrading its

system ' incorporate state - of- th -art advances. NIST is then the logical

host for the technical meeting. Ihe personnel involved with programming in

each laboratory are: D.T. Read, NIST; I. Hwang, MIT; R. Riddle, LLNL; and

J. Chan, LBL. A date has not yet been set, but all laboratories have agreed

in principal to a meeting. In addition to the above, more compact specimens

are requested from JAERI so that re -tests can be performed as necessary.

ASTM TECHNICAL COMMITTEE MEETINGS

Temperature Effects Subcommittee (ASTM E28.10).

This committee is processing our 4-K tensile standard. Mr. Thomas

Wilmink (Inco Alloys International, Huntington, West Virginia, 25720;

telephone: 304-526-5690) has replaced R. McDemus as chairman.

Every standard must pass three ballots: the subcommittee, main

committee, and society wide ballots. At this meeting the subcommittee

ballot results for the tensile standard were made known. Our standard

passed with 5 affirmatives, 5 abstentions, and 5 not returned. The standard

now graduates to the second step, which is main committee ballot.

Regarding responses to our standard there were no negative criticisms,

but a professional outside the subcommittee submitted some useful

suggestions in writing. Also, at the 1988 U.S. — Japan workshop we agreed

to define and add a new parameter, o^, the stress at the first serration.

The subcommittee faced a choice of revising at this point or waiting for

revisions until after the standard went through all the voting stages. At

first it was decided to proceed without a revision, for the sake of

expediency. Later, the chairman has reversed that decision because he

discovered that a revision could be made without any substantial delay in

processing at ASTM. Therefore I have agreed to provide a newly revised

draft before the end of March. The subcommittee will have the revised draft

at the next technical meeting (May 22, Kansas City). That draft will then

be submitted to the main committee ballot in September, and any negatives

220
^SP«!*5sr

will be resolved at the technical meeting in November, 1989. According to

this timetable the standard should be accepted in 1990.

Regarding developments on the other side, I learned recently from

H. Nakajima that standardization of the tensile test is also proceeding in

Japan. The original draft standard was revised slightly to simplify the

requirements for alignment, and it is being processed as a Japanese

Industrial Standard by a committee with Dr. H. Takahashi serving as

chairman.

Fracture Initiation Concepts Task Group (ASTM E24.08.04).

This committee is processing our 4 K fracture toughness standard, among

numerous other activities. At this meeting there were seven technical

presentations, including my own summary of recent developments in 4-K

testing, a talk highlighting data from the U.S. — Japan workshop (May,

1988).

In the presentation and the ensuing discussion, I emphasized that the

austenitic steels at 4 K do not behave exactly like ferritic steels at room

temperature. Evidence presented at the U.S. — Japan international workshop

had indicated that the blunting behavior, regression line formalisms, and

thickness requirements are different. In this way I reasoned to accept all

of those modifications that the American and Japanese scientists agreed to

at the international workshop.

Drs . Joyce and Hackett co-chair this task group, but Ed Wessel

exercises main committee supervision. At Vessel's suggestion it was agreed

to revise the format of our standard; instead of a separate new document,

our standard will be re-written and included as an Annex to the existing

method E813-87. Joyce is still somewhat concerned that our proposed

modifications for cryogenic austenitic steels will conflict with the main

text of E813-87. Therefore I must discuss and resolve the technical points

with him by personal correspondence.

Regarding a schedule for progress, I agreed to write and distribute the

Annex by April 1, 1989. This Annex must be acceptable to the co-chairmen

and will be formally presented to the task group at the next technical

meeting (May 22, 1989 in Kansas City). Balloting of the Annex should begin

then.

According to the above, there must be a major change in our 4-K

fracture standard. The present draft combines Kj c and Jj c testing, whereas

the Annex will cover only Jj c testing at 4 K. NIST agreed to this change

for several reasons. First, we believe that the existing Kj c standard (ASTM

Methods E399-83), although it does not reference cryogenic apparatus, is

fully satisfactory for 4-K testing in regard to its technical requirements

(loading rate, precracking guidelines, etc). Also, the combination of test

methods is already being pursued by another committee (ASTM E24.08.01).

That group is working to establish a test procedure consisting of one large

document covering and combining Kj c , Kr, Jj c , Jr# anc* CTOD tests. They are

exploring modifications of all test procedures in an attempt to achieve

conformity between linear elastic, elastic-plastic, and fully plastic tests.

Therefore it seems better for us to complete the 4-K Jjc test Annex now and

benefit from the ASTM work in the future.

221
Composite -Related Meetings

Committee D-30 (High Modulus Fibers and Their Composites) also held

several meetings in Atlanta, but they conflicted with the metals meetings.

I missed the Shear Testing Task Group (D30.03.01) and part of the Study

Group on Compression Testing (part of D30.90). However, I attended the main

committee meeting D-30 which provides an overview of all activities.

The status of standards for composites testing at room temperature is

complicated. For example, ASTM now has three compression test standards and

three more are in process. The question is: to what extent can we develop

a single cryogenic test method that can address the various different design

needs? One approach is to review the existing tests and relevant

literature, and select a method with the features most relevant for

cryogenic testing.

Committee D-20 (Plastics and Reinforced Plastics) did not meet in

Atlanta, and generally there is little interaction between D-20 and D-30.

Our main interest is glass reinforced plastics, but the methods of D-30 may

prove better for us than those of D-20 because D-30 pays more attention to

the problems of stress concentration effects which are of greater concern in

cryogenic tests.

CONCLUSION

The developments in cryogenic materials test standards may be

summarized as follows: The 4-K tensile standard is well into the balloting

process and may be published in 1990. The 4-K fracture standard has been

reviewed by ASTM, but it must be re -formatted before balloting begins. I

will send the revised drafts of the tensile and fracture standards to ASTM

and to Japan according to the schedules indicated above. At present, there

are no cryogenic standards for composites; we can begin by formulating a

strategy for compression and shear tests based on what can be gained from

experience at room temperature. A special session will be created at ICMC

in July to gain background for standardization activities for composites.

222
UNITED STATEB DEPARTMENT OF COMMERCE

National Inatitufea of B«anclarda and Technology

[formaHy National Bureau of 8tandarda]

325 Broadway

Boulder, Colorado B0303-332B

Reply co ere accencion of:

March 21, 1989

MEMORANDUM FOR: Dr. H.I. McHenry

Chief, Fracture and Deformation Division

From: R.L. Tobler

Fracture and Deformation Division

Subject: Trip Report - ASTM D30 Symposium and technical committee

meetings related to development of standards for mechanical

property tests of composites (Charlotte, NC, March 14-16,

1989)

I attended these meetings to gain contacts and information on the

current activities of ASTM and other groups interested in the development of

test standards for composites. Our goals for the DOE program include

eventual development of test standards for composites at cryogenic

temperatures

.

I. SYMPOSIUM ON AUTOMOTIVE AND INDUSTRIAL COMPOSITES

This symposium ran into problems as 4 out of 8 papers were cancelled.

However, the cancellations created extra time for informative discussion.

The issues in this field unfold as follows:

• Accelerated testing - there is a need for tests enabling long term

effects to be predicted from short-term measurements.

• Combined effects - there is a need for tests enabling the

prediction of long-term durability of composites subject to

combined effects of moisture, temperature, and stress.

• Fiber distribution effects - there is a need to account for fiber

orientation distribution variations created by processing

operations

In automotive applications, the use of composites results in improved

quality, durability, damage tolerance, and reduced costs or weight savings.

Composites are already used in external body panels as well as some primary

components: glass-epoxy is used in leaf springs and glass/graphite hybrid

composites are used in drive shafts. Uses are expected to expand in the

future

223
An Automotive Composites Consortium (ACC) has been formed to improve

America's international competitiveness by sponsoring general research of

interest to the three participating manufacturers (Ford, GM, Chrysler). The

ACC consists of a board of directors and three task groups: Materials,

Processing, and Crash Energy. Their goal is to develop a textbook of

composites technology including test methods and databases useful to all

members and serving as a guide for their materials suppliers.

ACC's approach to standards is to adopt and adapt existing documents

(ASTM, SACMA, and MIL standards), or create new standards as needed. At

present structural components and processing are being emphasized. The

automotive industry differs from the aerospace industry in that low cost,

quick, and more specific tests are needed. Material variability is more of

a problem because production techniques tend to create distributions of

fiber orientations in contrast to the more highly oriented and well

specified materials use in aerospace.

II. COMPRESSION TESTING STUDY GROUP

Various aspects of compression testing were considered here in three

presentations. Dale Wilson's description of SACMA and its cooperative

interaction with ASTM was by far the most significant item.

SACMA (Suppliers of Advanced Composite Materials Association) is a

trade association strongly pushing the development of test standards for

commercial purposes. They have already drafted eleven standards and are

submitting them for processing through appropriate committees in ASTM.

Wilson (Celion Carbon Fibers, Inc), an invited speaker for our next

ICMC meeting, and active in both SACMA and ASTM, presented three of SACMA'

s

draft standards for oriented fiber-resin composites to this D30 meeting and

proposed prompt execution of round robin test programs supporting each

standard.

Wilson described the materials to be tested, the fixtures to be used,

and the test methodologies to be evaluated (compression, compression after

impact, and open hole compression). Each methodology would be evaluated

using 3-5 materials, and 5 tests per material. The study group voted in

favor of his proposal, appointed specialists to review the draft documents,

and agreed to respond formally to SACMA within one month. As a result, test

results could be available at the next study group meeting (Orlando FL,

November, 1989). Ten labs are needed for each type of test, and I placed

NIST on the mailing list for possible participation.

III. AUTOMOTIVE/INDUSTRIAL SHEAR TESTING

There were three items on the agenda for this meeting. These included

comments on the automotive/industrial composites symposium (already

discussed), an update on the task group's proposed V-notch shear test

standard, and a presentation of a new method for in-plane shear measurements

using a compact double -notched specimen.

224
As currently written, the proposed V-notch shear test standard

recognizes both the Iosipescu and bend specimen geometries. This standard

received some negative votes in recent balloting, and the equivalence of the

two specimens became questionable because recent round robin results show

some outliers for the bend specimen. Consequently, the task group at this

meeting decided to omit the bend specimen from the document, proceed with

standardization using only the Iosipescu specimen, and introduce the bend

specimen later after additional study.

Post and Ifju (University of Virginia, Blacksburg) described a new

specimen for in-plane shear measurements and presented preliminary

experimental work using cross-ply specimens of graphite -epoxy. They find

the new specimen geometry to be eminently suitable for routine

determinations of stress -strain properties in shear, and for shear strength

measurements at ambient temperatures.

The new specimen, a compact double -notched geometry, was 1.5 in long,

0.75 in wide, and 0.25 in thick, with opposing notches. Post claims from

moire analysis that this specimen offers a state of nearly pure and uniform

shear in the region between the notch tips, and that it may be superior in

this respect to the popular Iosipescu specimen. The compact specimen is

small, the fixtures are simple and inexpensive, and the tests are easy to

perform. I assume the specimen can be scaled up or down in size a little

more. Therefore, from our own perspective, it appears that this compact

specimen may be the ideal choice for shear tests of irradiated composite

materials at 4 K.

IV. CONCLUSION

The choice of specimen design is one of the critical steps in

developing cryogenic test standards for shear and compression testing.

Based on the information gained at these meetings , it appears that the

problem of specimen geometry for 4-K shear testing of irradiated composites

may be solved by using the new double -notched compact specimen. Laboratory

tests of this specimen at cryogenic temperatures are needed to verify

applicability and optimum specimen dimensions.

225
,
CHARPY TESTS NEAR ABSOLUTE ZERO

R.L. Tobler and R.P. Reed

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado 80303

Because the specimen warms up as it is transferred to the

test machine and fractured, valid impact tests of alloys at 4 K

can not be performed according to the standard procedure outlined

in ASTM Methods E 23-86. This paper describes the problems of

Charpy testing at extreme cryogenic temperatures, especially 4 K,

including alternative methods of testing and calibration. Ve

review previous experience, including observations of specimen

temperatures during various phases of testing, and correlations of

the Charpy impact energy (CVN) with quantitative toughness

parameters such as Jjc or Kj c

.

During Charpy tests at impact strain rates the temperature of

austenitic steel specimens initially at a nominal temperature of

4 K rises over 100 K, and correlations between CVN and Kjc or Jj c do not necessarily show the desired 1:1 correspondence. Fracture

does not occur at the intended temperature, or even at a constant

temperature, since materials with different work hardening rates

show different adiabatic heating behavior. Because of the

temperature rise variability and data scatter in the measurements

and property correlations, we conclude that it is not possible to

accurately estimate the fracture toughness of ductile steels, or

rank them exactly, using Charpy tests at 4 K.

INTRODUCTION

Simple methods of fracture toughness evaluation are always desirable,

especially at cryogenic temperatures where the tests are more expensive.

Traditionally, the Charpy test has been used to evaluate the toughness of

materials, especially ferritic steels, at temperatures as low as 77 K. For

aerospace applications the temperature of interest is 20 K, and for superconducting machinery applications it is still lower at 4 K. Researchers

have performed Charpy tests at these lower temperatures, but experience

shows that special procedures and precautions regarding data interpretation

are necessary when the test temperature approaches absolute zero.

227
The characterization of structural alloys for large superconducting

magnets in fusion energy machines requires the evaluation of tough austenitic alloys at 4 K. Computerized J-integral fracture tests at slow

loading rates provide a reliable quantitative measure of fracture toughness

at 4 K; simpler Charpy tests have been used for screening purposes to

develop new cryogenic alloys, and are sometimes proposed for quality control

or acceptance tests. As argued in this paper, however, the meaning of

impact test data at temperatures near absolute zero is questionable because

of problems associated with temperature control.

The existing standard [1] does not mention it, but Charpy testing at

4 K requires departures from the conventional procedure because the specimen

after being cooled to 4 K cannot be transferred through air to the test

machine in the routine way. Instead, various methods of insulating the

specimen on the anvil are used. The insulation is fractured with the

specimen, necessitating a nonstandard calibration correction. Although

specimen heating during transfer can be eliminated that way, the severe

adiabatic heating of the specimen upon impact inevitably complicates any

interpretation of the results. Adiabatic heating occurs in mechanical tests

at any test temperature, including room temperature, but is much worse at

extreme cryogenic temperatures because of the very low specific heats of the

test materials.

Most laboratories that conduct impact tests rely on swinging pendulum

type machines. A few labs have special facilities for nonstandardized

tests, but little information on these other methods is available in the

literature. As illustrated photographically by Novikov [2] and schematically by Verkin [3], a pneumatic device is used in the Soviet Union for

impact tests at 4 K. Their device consists of an enclosed chamber with a

vertically aligned and pneumatically driven hammer. Drop-weight impact

tests [4] and instrumented impact test [5] at 4 K have also been conducted

on a limited scale, but only conventional swinging pendulum tests are

considered in this paper.

PROCEDURE

CQTtveutipnftl Testing

Conventional Charpy testing is conducted according to the ASTM Standard

Methods for Notched Bar Impact Testing of Metallic Materials, Designation

E 23-86 [1].- Within 5s, a 10 x 10 x 55 mm notched bar specimen is quickly

removed from a cooling bath, positioned on the anvil of an impact test

machine, and fractured. The absorbed energy and the lateral expansion of

the specimen at the notch plane are then measured to characterize fracture

resistance. This procedure is quite satisfactory at temperatures down to

77 K, especially for the traditional characterization of ferritic steels,

but for austenitic alloys at still lower temperatures it proves inadequate.

Three problems are encountered: the temperature of the test specimen rises

as it is transferred to the anvil, the specimen is susceptible to adiabatic

heating as it fractures at impact strain rates, and the data scatter for

replicate tests increases at cryogenic test temperatures.

228
Cryogenic Tasting

The most common temperatures in cryogenic research are the fixed bath

temperatures of the liquid cryogens: 77 K (nitrogen), 20 K (hydrogen), and

4 K (helium) . The principal means of verifying the adequacy of tests at

these temperatures has been to instrument calibration specimens with

thermocouples and record the specimen temperature throughout the various

stages of testing.

Depending on purpose, the thermocouples can be attached to the specimen

surface, or implanted in the center of the thickness by means of drilled

holes. Chrome 1-versus-Go Id + 0.07% iron thermocouples are preferred for

maximum sensitivity below 20 K. Experimental observations confirmed that

for specimens initially cooled to 77 K there is no appreciable temperature

rise within the 5 s that it takes to transfer and fracture the specimen

according to the ASTM procedure [6]On the other hand, data for specimens

cooled to 20 or 4 K show that the temperature begins to rise immediately

after the specimen is removed from its Dewar [7,8]. Therefore different

procedures are needed, depending on the intended initial test temperature:

tests at 77 K or above can be conducted by the established procedure

specified in ASTM Methods E 23, but to conduct tests at 20 or 4 K we must

resort to special techniques. A number of approaches are described in the

literature, but none are specified in the current standard.

In the quest to obtain satisfactory results at 4 K, almost every conceivable approach has been tried, including direct transfer of the specimen

from the cryogenic both through air to the anvil of the test machine,

encapsulation of the specimen, and encapsulation of the entire test machine.

Each method offers specific advantages and disadvantages, and none has come

to predominate. A number of variations of each approach are considered below.

Direct Transfer

Several researchers have tried to apply ASTM Methods E 23 while making

no special provision for extreme cryogenic test temperatures. Their efforts

show only that the conventional procedure fails.

For example, Lee and Dew Hughes [9] reported Charpy data for stainless

steel castings at several temperatures as low as 77 K but they were forced

to abandon Charpy testing at 4 K because "the problems of temperature

control and scatter were unacceptable”. To document the problem of temperature control quantitatively , Dobson and Johnson [8] recorded the thermal

histories of AISI 304 specimens containing embedded thermocouples. Using

pre-cooled tongs, they transferred bare specimens through air as quickly as

possible but the specimen temperatures increased to 67-71 K before the

specimens could be impacted. Yoshimura et al. [10] did better by covering

their specimens during transfer; still, the temperature of their specimens

rose to about 15 K before the hammer could strike them.

A rise to 15 K or more is of course unacceptable, since a temperature

accurate to within one kelvin is required by the standard. Therefore significant efforts have been made to insulate the specimen, and sometimes the

entire machine, to achieve a fixed initial test temperature. Several

methods are reviewed below.

229
Boating. Boxing, and Encapsulation

A paper boat method for tests at 20 K was first developed at Ohio State

University [7]. As described by Zambro and Fontana [6], a standard Charpy

specimen is glued to the bottom of a paper boat that has a perforated cover

and is suspended by threads to paper tabs used for handling. After cooling

in liquid hydrogen, the specimen is lifted out of its dewar and transferred

through air to the anvil of the machine. During transfer, the container

retains sufficient liquid to forestall specimen heating. Once the specimen

is on the anvil, the threads are cut and the hammer released. A clever

variation of the boating method can be used for tests at 5 or 6 K, as

described by Long [11].

Although the paper boat method and its variants permit tests at 20 or

lower, the method is disadvantageous in that hydrogen is hazardous, special

precautions are necessary to avoid explosions. Therefore Jin et al. [12]

developed a test using liquid helium and a plastic insulating box for

specimens tested below 6 K.

To improvement on the method described by Jin et al., Ogata et al.

[13,14] sought simpler types of encapsulation. In their system, the plastic

box is replaced by a capsule made of grooved polystyrene foam which achieves

better cooling efficiency and minimizes the calibration error.

Another recently developed technique, widely used in Japan, is the

glass Dewar method [15]. The specimen is placed in an (open or closed)

mouth vacuum- insulated glass tube capable of retaining liquid helium for

several minutes. The tube containing the specimen is then placed on the

anvil and fractured. This method assures an initial test temperature of

4.2 K, but the CVN measurement must be corrected to account for the significant energy absorbed in fracturing the glass container. Two methods of

calibration correction have been suggested, as discussed in the next section.

In other work, the test machine as well as the specimen was encapsulated [7,16]. In general, the unique advantages of those methods are

outweighed by the complexity of the apparatus required, and the methods

haven't been used by anyone except the original experimenters.

Calibration

Tests of encapsulated specimens require a nonstandard calibration

correction to account for the energy absorbed in fracturing the insulation

or container. Inevitably, assumptions are involved in arriving at a

correction factor.

Usually, tests of encapsulated or insulated specimens are calibrated

simply by subtracting the energy required to break the capsule from the

total energy. The correction factor is deduced by comparing the results of

tests performed on empty capsules versus the results for capsules with

specimens. The comparison is sometimes made at 273 K [15], but usually at

77 K [17]. The calibration correction factor so determined is then used

also at 4 K, assuming that the effect of the test temperature on the

calibration factor is negligable. The underlying assumption, however, that

the energy to break the capsule with a specimen is equal to the sum of the

230
energies required to break an empty capsule plus that to fracture a bare

specimen, has not yet been experimentally verified.

Typical calibration factors for the various methods are listed in

Table 1. The paper boat method has the advantage of the smallest correction

factor. For the glass Dewar encapsulation technique, the correction factor

is quite high and, for low toughness materials, it would amount to a

significant fraction of the total impact energy. The glass Dewar method is

therefore not suited for the low toughness alloys.

Another calibration technique is possible, but less commonly used. In

Charpy testing it is customary to measure both the energy absorbed on impact

(CVN) and the lateral expansion (LE) . A number of researchers have shown

that CVN and LE are approximately linearly related for certain conditions,

and, as described by Mori and Kuroda [18], this relationship provides a

basis for a second calibration procedure for tests at 4 K.

Adiabatic Heating

While specimen warming during transfer can be eliminated by boating,

boxing, or encapsulation, the heat generated in the specimen during deformation and fracture at impact loading cannot be avoided.

Adiabatic heating depends on material and test variables, including the

total quantity of plastic work consumed in fracturing the specimen, the rate

of straining (heat production), the specific heat of the test material, and

the efficiency of the cooling medium. Austenitic steels are more susceptible than other steels because they are usually tougher and have higher

specific heats. In fact, adiabatic heating produces a measurable temperature rise in austenitic steel specimens during conventional tensile tests at

room temperature [19]. The problem is worse at impact strain rates, and

still worse at cryogenic test temperatures where the specific heats of the

test materials are drastically reduced. For example, the specific heat of

AISI 310S steel decreases by several orders of magnitude for a temperature

change between 298 and 4 K.

Table 1. Typical calibration correction factors for Charpy tests methods

at extreme cryogenic temperatures.

Method and type

of insulation

Test temperature

K

Calibration

J

Reference

paper boat 20 0.3 [6]

polystyrene foam 5 1.5 [13]

plastic box 5-6 4.9 [12]

glass dewar 4.2 33 [15]

231
Adiabatic heating during fracture can be studied using specimens

instrumented with thermocouples. In their impact tests of AISI 304 steel,

Dobson and Johnson [8] recorded temperatures as high as 150 K. In future

work, we intend to measure the temperature rises and compare them with

numerical calculations based on the conversion of plastic work to heat.

Data Trends

Typically, Charpy data for ferritic steels over a range of test

temperatures show an abrupt decrease of energy at the transition temperature, whereas data for austenitic steels and alloys show a gradual increase

or decrease of impact energy with test temperature. The same general trends

have been observed in the extreme cryogenic range. For example, data for a

9% Ni steel and a high Mn austenitic steel were compared by [20]

.

The

austenitic steel showed a gradual decrease of CVN at five temperatures

between 4 and 295 K, whereas the ferritic steel showed a typical transitional behavior. There is one noteworthy fee ure regarding ferritic steel

transions: most mild steels have a transition temperature in the range

275—375 K, but in higher Ni steels the transition temperature is suppressed

to 77 K or lower because Ni beneficially affects the resistance to cleavage.

Data Scatter

Data scatter is a principal shortcoming of the Charpy test, and the old

problem is magnified when Charpy tests are performed at temperatures

approaching absolute zero. During the time it is being transferred, a bare

test specimen warms continuously. Therefore the actual specimen temperature

at the time of testing must vary with the transfer time, and this contributes to the data scatter of results.

Also, the scatter in data at extreme cryogenic temperatures increases

relative to the level observed at room temperature because metallurgical and

mechanical test variables have a greater influence on material properties at

cryogenic temperatures. A third factor contributing to the increased

scatter in Charpy measurements at 4 K is the added experimental uncertainty

associated with the calibration correction factors required at 4 K.

Pet? correlations

In a large program sponsored by JAERI in cooperation with several

Japanese steel companies, CVN data were used to screen newly developed

cryogenic steels [20-21]. In this program, ninety alloy compositions were

Charpy tested at 4 K, using the glass Dewar method. Alloys having CVN

greater than 80 J were selected for further study. The selected alloys were

then produced in greater quantities and subjected to quantitative Jjc tests.

A 4-K fracture toughness of Kjc — 200 MPa* nr* or more was required to satisfy

design constraints. In this way, six new alloy steel compositions were

developed for superconducting machinery applications. After the Jjc values

were measured, it was possible to compare the CVN and Jjc data for a number

of alloys. The results, however, did not show a perfect 1:1 correlation

[28, 29].

232
DISCUSSION

Ve are sometimes tempted to avoid the extra effort involved in testing

at 4 K by using 77 or 20 K data as a basis for making decisions about 4 K

applications. Mazandarany et al. [30], in a presention of data for an AISI

316LN base metal and several of its velds, demonstrated that this approach

is inherently unreliable. They measured CVN at 77 K as well as Jjc at 4 K,

but found no general correlation between the two parameters. Their data

show that the problems of fracture characterization at 4 K cannot be

resolved by Charpy tests at other temperatures. This is because the

deformation mechanism at 4 K, is unique from that at 77 K or higher temperatures. At temperatures from ambient to 77 K, the load- displacement curves

are regular, shoving slow stable plastic deformation. At 4 K, however, the

curves show instabilities which are due to increments of unstable (freerunning) plastic deformations, each accompanied by temporary thermal spikes.

While general correlations between Charpy data and fracture toughness

have been demonstated at ambient temperature, there are several reasons to

expect that CVN data should not correlate exactly with Kjc data at 4 K.

First, the two tests do not measure exactly the same properties of materials. The Charpy test involves crack initiation from a notch, as well as

crack propagation, whereas Kjc and Jjc pertain only to the initiation of

fracture from a pre-existing crack.

Second, the tests are conducted at different strain rates, and the

discontinuous plastic flow (serrated yielding) process that is a prominent

feature of the 4-K tests is very sensitive to strain rate. Especially at

nominal strain rates higher than about 10' 3

, the discontinuous deformation

behavior and adiabatic temperature rise for most austenitic steels changes

drastically [31].

Charpy impact data are used mainly in mechanical testing of steel

products, in research, and in procurement specifications. As stated in the

standard [1], the Charpy test is especially practical for tests of ferritic

steels because it enables the ductile- to-brittle transition temperature to

be determined; it is less informative when applied to austenitic alloys or

face -centered cubic metals that typically do not exhibit transitional

behavior. On these grounds, ve conclude that the test is not ideally suited

for austenitic alloys, or for the temperature of interest. Since the

changes for austenitic alloys are not drastic, it is never clear to what

extent the CVN changes with test temperature are due to material behavior or

heating effects. Any test loses significance as the sensitivity of measurements approaches the property changes to be measured.

It seems that up to now the limitations of Charpy testing at liquid

helium temperature have not been widely appreciated outside of the relatively small community of cryogenic researchers. For example, regulatory

agencies today still require that helium storage tanks be designed according

to the ASHE boiler code which requires Charpy tests of materials of construction at 4 K. Although the value of Charpy tests at 4 K is questionable

on technical grounds, the tests are nevertheless conducted because the code

requires them. However, neither the code nor the Charpy test standard were

originally written with 4-K applications or experience in view, and as

emphasized in this paper, it is actually impossible to perform valid Charpy

233
?ugs&jr.-r*

tests at 4 K according to Methods E 23-86 as currently written, and as

required by regulatory agencies.

Based on the facts reviewed in this paper, it may be appropriate to

revise certain sections of Methods E 23 to recognize the problems of testing

at extreme cryogenic temperatures. Revisions may include a temperature

limitation for cryogenic testing to 77 or 20 K, or explicit recognition of

special techniques appropriate for testing at 20 or 4 K.

CONCLUSION

In this paper, we emphasize that it is impossible to perform Charpy

tests at 4 K according to the procedure outlined in ASTM Methods E 23.

After the specimen is removed from its dewar and before the hammer strikes

it, the specimen temperature is continuously rising.

Charpy testing at 20 or 4 K entails a departure from the standard

procedure of ASTM Methods E 23-86 because the specimen cannot be cooled and

subsequently transferred to the test machine in the routine way. This paper

describes the various alternative methods that have been used to perform

impact tests at 4 K. Cryogenic researchers always employ some nonstandard

procedure, and most prefer one of the various methods of insulating the

specimen on the anvil. The insulation is fractured along with the specimen,

requiring a nonstandard calibration correction. Each method has specific

advantages and disadvantages , and none has gained predominance

.

The Charpy test loses its simplicity in view of the extra efforts that

must be taken to obtain data at 4 K. Consequently, the value of the Charpy

test at 4 K is compromised. Perhaps some merit as a crude screening

parameter for identifying brittle materials may be retained, in that a low

CVN generally implies low toughness.

REFERENCES

[1] Standard Methods for Notched Bar Impact Testing of Metallic Materials,

Designation E 23-86, in 1988 Annual Book of ASTM Standards, Section 3,

Metals Test Methods and Analytical Procedures, Amer. Soc. Test.

Maters., Philadelphia, 1988, pp 191-206.

[2] Novikov, N.V. , Mechanical Property Measurement Techniques of Structural Materials at Cryogenic Temperatures, in Adv. Cryo. Eng., Vol.

22 . K.D. Timmerhaus, R.P. Reed, and A.F. Clark, Eds., Plenum, New

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237

WARM PRECRACKING AT 295 K AND ITS EFFECTS

ON THE 4-K TOUGHNESS OF TWO AUSTENITIC STAINLESS STEELS*

R.L. Tobler

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, CO 80303

M. Shimada

Iron and Steel Technology Center

Kobe Steel, Ltd.

Kobe, Japan

Experiments reported here show that some austenitic steels

are toughened by warm prestress. We demonstrate this for Fe-17Cr3Ni-13Mn-0. 33N steel specimens that were warm precracked at 295 K

and then fractured at 4 K where failure occurs by slip-band

cracking. On the other hand, the toughness of Fe-13Cr-5Ni-22Mn0.21N steel with a ductile fracture mechanism at 4 K was not

affected by similar warm precracking. The behaviors of austenitic

and ferritic steels are compared, and valid fatigue precracking

procedures for standard 4 K fracture tests are considered.

INTRODUCTION

Preloads applied to steel specimens or structures at a relatively high

temperature Tj can influence the fracture toughness subsequently measured at

a lower temperature, T2

. Either a single load or cyclic loading at T x may

increase the toughness at T2

. This, the warm prestress (WPS) effect, is

common in ferritic steels when Tj is a temperature above the ductile-tobrittle transition and T2 is below the transition temperature. As shown in

this paper, toughening from WPS also occurs in some austenitic steels if Tj

is room temperature and T2 is near absolute zero.

Theories about WPS were advanced by Chell, Haigh, and Vitek [1], and by

Curry [2]. Reviews were written by Nichols [3], and by Pickles and Cowan

[4]

Previously, WPS effects were observed only in steels that had ferritic

microstructures. Some examples include high-C spheroidized steel [1],

medium strength steels [5], C-Mn steels [6], 0.36% C steel [7], A 533 B-l

steel [8], and free machining steel [9].

239
Our interest in austenitic steels began with questions about fracture

test procedures at 4 K. Some investigators like fatigue -precrack their

specimens at room temperature, for convenience, trusting that the toughness

at cryogenic temperatures will not be affected. The supporting data are few

and restricted in scope to Fe-Cr-Ni steels with ductile failure mechanisms

at 4 K: in recent studies of the 300 series stainless steels, precracking at

76 K - 60 MPa -in 1 ) did not affect the toughness of AISI 310S at 4 K

[10], precracking at 293 or 77 K (1^** - 24 MPa-m

1*) did not affect the

toughness of AISI 316 LN at 4 K [11], and precracking at 295 or 77 K (25 s

Kjnax 3 43 MPa-in 1

) did not affect the toughness of SUS 304 steel at 4 K [12].

This paper contrasts two high -manganese austenitic stainless steels

which have very different 4-K fracture characteristics. The effect of

fatigue precracking at room temperature (RT) or liquid nitrogen temperature

(LNT) on the fracture toughness at liquid helium temperature (LHeT) is

investigated. We refer to room temperature fatigue precracking of 4-K test

specimens as warm precracking and consider it a form of WPS. We find a

toughening effect at 4 K after warm precracking for the relatively brittle

steel, but not for the tough steel.

MATERIALS

The two materials studied were a commercial heat of Fe-17Cr-3Ni-13Mn0.33N steel (designated 13Mn) and a laboratory heat of Fe-13Cr-5Ni-22Mn0.21N steel (designated 22Mn) . The mill chemical compositions in weight

percent are given in Table 1. Both materials have high manganese and

nitrogen contents, and both are annealed.

The 13Mn steel is a high strength steel, selected because it has a

cleavage- like failure mechanism at 4 K, and experience indicates that steels

cleave are often susceptible to WPS effects. A 25.4 mm thick plate was

obtained in the commercially annealed condition.

The 22Mn steel is a newly developed cryogenic steel which has high

strength and toughness at 4 K [13]. This steel is a candidate for use in

magnets for fusion energy devices at 4 K, and it represents a typical case

of ductile fracture at 4 K. The test material was a 90 kg laboratory heat,

vacuum- induction melted, forged (1200°C), hot-rolled (28 mm plate), solution-treated (1050°C), and water -quenched.

Table 1. Chemical compositions of the test materials.

Material Fe Cr Ni Mn N C Si S P

13Mn bal. 17.3 3.42 12.66 0.33 0.05 0.50 0.003 0.027

22Mn bal. 12.92 4.93 21.9 0.212 0.02 0.50 0.004 0.003

240
PROCEDURE

The 13Mn and 22Mn steels were tested at two different laboratories.

Depending on the geographical elevation of the test laboratories, nitrogen

bolls at 76 or 77 K, and helium boils at 4 or 4.2 K, owing to the different

atmospheric pressures. In the text, these test temperatures are designated

by nominal values: 77 and 4 K, respectively.

The conventional tensile properties were measured using transversely

oriented cylindrical specimens of diameter D and length of reduced section

LRS. For 13Mn steel, D - 6.35 mm and LRS - 38.1 mm. For 22Mn steel, D -

7 mm and LRS — 39 mm. Table 2 shows the average tensile property measurements for two tests per condition.

Fracture toughness was measured using T—L oriented compact specimens

25.4 mm thick (B) and 50.8 mm wide (W)

The procedure followed a fatigueload, unload, cool to 4 K, and re -load to fracture sequence. A series of

specimens of each steel were precracked at RT or LNT and then fractured at 4

K (Tj > T2 - 4 K) by standard methods [14-16]. The results were compared

with data for control specimens that were precracked and fractured at LHeT

(T x — T2 - 4 K) . Post-test examinations included scanning electron microscope (SEM) observations of the fractured surfaces at locations near the

fatigue crack tip.

The 13Mn steel specimens were evaluated using ASTM Method E 399-83 [14]

to measure the critical stress intensity factor, Kj c

. The specimens were

precracked in fatigue automatically [17]. Load cycling was performed with a

sinusoidal waveform in load control at a frequency (f) of 30 Hz while the

crack length (a) was monitored by compliance. The fatigue -stress- intensity

factor range (AK) varied from specimen to specimen, but was maintained

constant during each test by periodic load reductions as a function of crack

growth. After precracking, the specimens were fractured by a single loading

at a crosshead displacement rate of 0.03 mm/s.

Table 2. Tensile properties of the tested steels (data are averages of two

tests per temperature at strain rate of about 2 x 10’ 4 s’ 1 ).

Material

Test

Temperature

(K)

Yield

Strength

(MPa)

Ultimate

Strength

(MPa)

Elongation

(%)

Red. of

Area

(%)

13Mn 295 401 786 59

76 1019 1489 20 - 4 1401 1733 7 -

22Mn 298 306 640 78 75

77 831 1315 59 43

4.2 1188 1565 38 42

241
The high toughness 22Mn steel was evaluated using Jxc measurements

,

from which Kjc

(J) values were derived. The specimens were precracked at

room or cryogenic temperatures at f - 5, 10, or 20 Hz. For this steel the

fatigue loads applied to each specimen were held constant, so the maximum

stress intensity factor in fatigue (Kmay) increased with crack length.

After precracking the specimens to similar relative crack lengths (a/W -

0.62) using different final AK, we fractured them using a computer-aided

unloading compliance apparatus [18] and a methodology that defines Jjc at

the intersection point of the blunting and regression lines [15, 18]. Systematically higher (10-15%) Jjc values would have been obtained using the

revised methodology [16], but that would not change the conclusions of this

study.

For both steels, the fatigue load ratio (R) was 0.1. The magnitude of

fatigue loading may be specified by either AK or Kjj^ since

K “ (^min/^max) " ^^min/^max) "0.1, and (1) “ 0 . 9Kmfl y . (2)

RESULTS

Linear Elastic 13Mn Steel

Load-displacement records for two 13Mn steel compact specimens are

shown in Fig. 1. The loading response is linear elastic, as expected. Most

of these tests satisfied the criteria of Method E 399-83, except where the

fatigue load restrictions were intentionally violated to create VPS.

Table 3 lists the test results for three control specimens (T t - T 2 -

4 K) , plus results for eight more tests where the fatigue precracking and

fracture test temperatures were dissimilar. The average Kjc value for

control specimens is 83 MPa'm*1

, with a spread of values between 81 and 86

MPa-m**. Higher toughness is observed for most of the warm precracked

specimens, but this depends on the prestress temperature T x and the fatigue

stress

.

The results for the 13Mn steel are plotted in Fig. 2 and summarized as

follows. If precracking is performed at RT (Tj - 295 K) , a VPS effect is

observed, and the measured toughnesses can be much higher than Kjc

. In

Table 3, the Kq measurements that are affected by VPS are designated Ky.

The highest Ky in Table 3 is 114 MPa-m*5

, 40% higher than Kjc

. Still higher

Ky could have been measured if the precracking loads were higher. As shown

in Fig. 2, Ky decreases and becomes negligible at or below K^y * 17 MPa*m . This value, for 13Mn steel, may be a threshold. Actually, we did not

demonstrate the existence of a threshold since we obtained no data below 17

MPa-m*1 (more than a million fatigue cycles per specimen would have been

necessary); however, thresholds are observed in some ferritic steels [3, 6]

and thresholds might be expected from the plastic zone size rationale

discussed below.

242
50

Figure 1. Load displacement curves for compact specimens of the 13Mn steel

tested at 4 K: (a) normal loading behavior, and (b) abnormal

behavior caused by inadequate precracking procedure (Kq^'0.6

Kic>.

Table 3. Precracldng and 4-K toughness results for 13Mn steel.

Specimen

No.

Precrack

Temperature

K

AK

MPa*in* MPa*nr MRmb* MPa*nr % KIc . MPa-in*

4 295 15.5 17 84.3 • « . 84.3

10 295 22.5 25 89.2 89.2 —

1 295 30.0 33 92.7 92.7 —

2 295 40.0 44 98.0 98.0 —

3 295 60.0 67 114.4 114.4 —

8 76 27.0 30 81.6 ... 81.6

12 76 35.0 39 82.0 — 82.0

9 76 54.0 60 80.7 — —

7 4 25.0 28 82.6 ... 82.6

5 4 30.0 33 81.1 — 81.1*

6 4 30.0 33 85.5 ... 85.5

243
T

Kmax* MPa-m1 '2

Figure 2. Fracture toughness of 13Mn steel at 4 K, showing the effects of

fatigue precracking conditions.

Some specimens were precracked at LNT (T x - 76 K) , but this did not

toughen the 13Mn steel at 4 K. Two specimens were precracked at 76 K, at

moderate AK, and their toughnesses were equivalent to Kjc for control

specimens. A third specimen was precracked at higher loads (K^y - 60

MPa*m**) and the 4-K toughness that was in fact slightly lower than Kjc. As

shown in Fig. 1(b), the test record in this case displays abnormal early

pop- ins, attributable to a violation of the precracking rule, K^y £ 0.6Kj c (see Discussion)

.

Fully Plastic 22Mn Steel

Tests of the 22Mn steel revealed no WPS effects for any of the conditions investigated in this study. A typical 4-K load-displacement record is

shown in Fig. 3, and typical 4-K J—resistance curves are shown in Fig. 4;

the Kjc estimates as a function of T x are shown in Fig. 5, and the fatigue

precracking histories for all specimens are listed in Table 4.

244
LOAD DISPLACEMENT, S, mm

Figure 3. A load—displacement curve for a compact specimen of 22Mn steel.

CRACK EXTENSION, Aa, mm

Figure 4. Resistance curves for 22Mn specimens precracked 298, 77, or 4 K,

245
Table 4. Precracking conditions and 4-K toughness results for 22Mn steel.

Specimen

No.

Temperature

K

AK

MPa-m1* Frequency

Hz

Crack Length

mm

Kic

(J)

MPa -mb

22Mn-l 298 20.4 20 31.9 186

22Mn-2 298 35.2 20 31.0 185

22Mn-3 298 33.5 20 31.5 a

22Mn-4 298 41.5 20 31.2 204

22Mn-5 298 45.9 20 32.7 187

22Mn-6 298 32.2 20 31.4 209

22Mn-7 77 30.1 20 31.1 197

22Mn-8 77 41.0 20 31.9 190

22Mn-9 77 56.3 10 31.6 191

22Mn- 10 77 56.6 10 31.1 a

22Mn- 11 4 40.1 10 30.7 185

22Mn- 12 4 30.6 10 30.6 199

22Mn- 13 4 36.4 10 30.7 196

22Mn-14 4 59.3 5 31.9 208

22Mn- 15 4 52.3 5 31.3 189

Note: aoperator error terminated this test;

The load—displacement records (Fig. 3) for 22Mn steel are serrated,

owing to repeated bursts of unstable plastic flow with sudden load drops and

arrests [10]. Serration size increases towards the end of the test. Such

behavior is typical of ductile steels tested in liquid helium at slow

testing speeds. Our 22Mn steel specimens were nearly identical in crack

size, so the serrations occurred at reproducible loads and displacements.

Consequently, the J—Aa data distributions within the exclusion intervals of

the various resistance curves are nearly identical, independent of precracking temperature

.

The fracture toughness data for 22Mn steel are plotted in Fig. 5.

Within the range of specimen- to- specimen variability, the toughness Kjc

(J)

at 4 K is 195 ± 7%, independent of the precracking conditions as far as they

were explored in this study. There are no WPS effects, either on the resistance curve regression lines or Jjc values for this material.

Fractographv

For both steels, SEM examinations of the fracture surfaces at locations

near the fatigue crack tip failed to reveal any obvious difference in

fracture morphology at T2 as a function of the precracking temperature (295,

76, or 4 K)

.

Representative fractographs are shown in Figs. 6 and 7.

For 13Mn steel, the 4-K fracture surfaces consist of transgranular

facets on {111} planes which are the active slip planes of the austenitic

structure. The 4-K failure mechanism is slip-band cracking, a form of

cleavage fracture [19]. Slip traces are visible on the transgranular

246
Figure 5. Fracture toughness values for 22Mn steel as a function of Kmax .

Figure 6. Fracture surface of the 13Mn austenitic stainless steel fractured

at 4 K.

247
Figure 7. Fracture surfaces of a 22Mn steel fractured at 4 K.

facets, as shown in Fig. 6. This implies some microplastic deformation

before fracture, even though the loading behavior (Fig. 1) is essentially

linear, and the overall fracture morphology (Fig. 7) is brittle.

In contrast, the high- toughness 22Mn steel affords a completely ductile

failure at 4 K. Figure 7 shows a specimen's fracture surface after Jj c testing at 4 K. The surface is covered with dimples from void coalescence.

Arrows in Fig. 7 indicate the direction of crack propagation. Figure 7(a)

is a macrophotograph of the specimen. Figure 7(b) is an SEM of the fracture

surface adjacent to the fatigue crack, and Fig. 7(c), an enlargement of the

same area, shows a distinctive feather- like dimple pattern. Figure 7(d)

shows the dimples at a distance 2 mm from the fatigue crack. As shown in

Figs. 7(b-d), and for all 22Mn steel specimens tested in this study, the

feather- like dimples exist only in a band adjacent to the fatigue crack tip.

The band is 200-500 pm wide, and is larger than the stretch zone.

248
Regarding the austenite stability of these alloys, X-rays of a previous

heat of 13Mn steel indicated no detectable austenitic phase transformation

(to bcc or hep martensite) , even after deformation and fracture at 4 K.

However, small amounts of transformation may have gone undetected because

the plastically deformed zone is extremely small. The present heat is quite

similar in composition to the heat that was X-rayed, so we assume it has a

similar level of austenite stability at 4 K. The 22Mn steel definitely

transformed to e martensite, but not to a

martensite, as determined by

radiography and electron microscopy and summarized in Table 5.

Table 5. Volume fractions of phases detected in 22Mn steel plastic zones.

Temperature Fatigue stress intensity austenite e martensite a' martensite

factor range, AK

K MPa * m^5 % % %

298 30 100 0 0

4 30-35 77 23 0

DISCUSSION

WPS has implications both for design and for materials testing. In some

designs with linear-elastic materials, WPS can be beneficial in that it

supports a strategy of intentional overstressing at T x to reduce the probability of fracture at T2 [4]. Our present interest, however, is valid test

procedure for fracture toughness measurements at 4 K.

There are good reasons for wanting to precrack at room temperature before

fracture testing at 4 K, Some laboratories are not equipped with 4-K fatigue

capabilities, and most desire to conserve liquid helium and reduce cryogenic

fluid costs. Since warm prestress effects may occur, it is necessary to

understand this behavior and consider acceptable fatigue precracking criteria

that will prevent warm precracking effects from interfering with standard

toughness measurements.

Comparison with Ferritic Steels

The behavior of the 13Mn austenitic steel of this study parallels

ferritic steel behavior in some respects. Figure 8 compares the 13Mn steel

and two ferritic steels which failed by cleavage at T2

. All three materials

are ductile at T x and brittle at T2 , and stress-controlled fracture mechanisms

(slip-band cracking or cleavage) operate at T 2

. All three steels were

preloaded at room temperature, but there is a minor difference in that the

13Mn steel was preloaded in fatigue; the ferritic steels were preloaded

monotonically

249
As Fig. 8(a) shows, the 13Mn steel is tougher than the ferritic steels.

One of the ferritic steels, a semi-killed Fe-0.19C-0.84Mn steel, was tested

at T2 - 77 K [6]; the other, a commercial free -machining Fe-0.17C-l.12Mn0.2Ni steel, was tested at T 2 — 123 K [9], Since the ferritic steels have

very low toughnesses (23 and 26 MPa-m^) at these temperatures the affected

toughness can be double, and the WPS effect seems relatively weak for the

13Mn steel in that a toughness increase of 40% was observed. The effects

are similar, however, if the data are normalized to account for the different Kp c values. The quantity Kyps - Ky

— Kj c is the increment of toughness added due to the WPS effect. As shown in Fig. 8(a), Kypg is a similar

function of Kmax for all three steels.

Significance of Plastic Zone Size

Toughness relates to the amount of energy absorbed in fracture which is

governed by the crack- tip behavior. Therefore WPS effects may be related to

plastic zone size effects, material composition and phase stability, and the

mechanical (cold-worked [20,21]) state of the material locally.

A factor to consider is the ratio of plastic zone sizes at T 1 and T 2

. Irwin [22] estimated the plastic zone radius for plane strain conditions as

Figure 8. Comparison of WPS effects in austenitic and ferritic steels:

(a) absolute data, and (b) normalized data.

250
This expression, although inexact [22] , is often used to estimate zone

sizes, and we use it here to estimate the magnitude of plastic deformation

associated with various precracking conditions and with fracture at 4 K.

Figure 9, then, indicates that there are large variations in the

fatigue plastic zone sizes developed at 295 and 4 K. In this figure, Ryf is

the radius of the plastic zone formed during the fatigue precracking, and

Ryp is the maximum radius of the plastic zone formed in control specimens

(those without warm prestress) at the point of fracture at 4 K. For 13Mn

steel at a typical precracking level (Kmax “ 33 MPa-rn^), the size of the

zone formed during fatigue at 295 K (Ryf) is 12 times larger than that for

identical loading at 4 K; it is also larger than Ryp_ The ratio Ryf/Ryf

appears to be significant. Figure 9 shows, for the steels tested, Ky = Kj c if this ratio approaches 1 or less. In other words, a WPS effect is likely

if the plastic zone formed in fatigue at T t exceeds that normally formed in

control specimens during fatigue and fracture at T

Significance of Fracture Mechanism

The literature suggests that WPS effects on toughness can be anticipated

in steels with stress -controlled (brittle) mechanisms at T 2 , and not in

steels with strain-controlled (ductile) failure mechanisms at T 2

.

In agreement, we find the effect in a 13Mn steel that is ductile at 295

K and brittle at 4 K, but not in the 22Mn steel nor other austenitic steels

that remain ductile at 4 K [10-12]. Therefore we expect to find WPS effects

in some other high strength austenitic steels, such as those of the Fe-CrNi-Mn-N family, which have high N contents and some proportion of crystallographic facetting at cryogenic temperatures [24-29]. The 22Mn steel,

which retains its dimpled- fracture mechanism at 4 K, shows no influence of

WPS at 4 K and behaves similarly to AISI 316 LN and SUS 304 [11, 12]. These

steels, in contrast to the 13Mn steel, are capable of extensive deformation

at extreme cryogenic temperatures and the deformation begun at T x is able to

continue at T 2

.

For 22Mn steel, feather- like dimples were observed on the fracture

surfaces. Presumably, this reflects the steel's microstructure and stability. Neither c nor a' martensite phases were induced by deformation at 295

K, but deformation at 4 K does induce some c (see Table 5). We suppose that

transformation at 4 K generates a layered structure of e and austenite

phases, that the € fractures first; then the austenite fractures in a more

ductile fashion, producing a textured surface and the overall feather- like

pattern. There is a change in fracture morphology, however, as the feathered pattern vanishes farther away from the fatigue crack (Fig. 7d) . This too

may be related to the transformation tendencies of 22Mn steel, since the e ceases to form at higher temperatures. From a previous study of fracture in

liquid helium, we know that, along with serration size, the internal

specimen temperature rises as testing progresses. Perhaps, after the crack

extends 200 to 500 /im, adiabatic heating associated with pronounced serrated

251
Figure 9. Fracture toughness at 4 K versus plastic zone size ratio.

yielding events is sufficient to suppress further transformation to c in

this alloy. The feathered pattern then vanishes because it derives from

fracture of a two-phase microstructure.

Unlike our high-Mn steels, many 300 series steels at 4 K readily

transform to martensitic products, including a phase. Martensitic transformations cause a volume expansion (2 or 3%) that depends on composition in

the 300 series steels. During fatigue or monotonic loading at cryogenic

temperatures, transformations occur in local plastic zones, but they occur

less readily, or not at all, at 295 K. Since loading at 295 K may affect

the phase transformation tendencies of steels like AISI 304 and AISI 316, it

might be possible to discover WPS effects in such alloys. Although previous

tests did not reveal any WPS effect in two of those alloys [11, 12], further

studies on very lean alloy compositions are recommended.

252
STANDARD CRYOGENIC TESTS

Test standards are designed to ensure valid fracture toughness measurements. For austenitic steels at 4 K, fatigue precracking restrictions are

needed to prevent the effects of loading history on fracture toughness

measurements. Guidelines should be simple and generally applicable. At

present, there are no fatigue precracking guidelines for 4-K Jj c testing,

and the published guidelines for Kj c tests have never been validated for

extreme cryogenic temperatures

Precracking at Dissimilar Temperatures

Method E 399-83 for Kj c measurements allows precracking at a higher

temperature T x , followed by fracture testing at a lower temperature T2 , but

only if Kmax at T x is not excessive. Section A2.4.4 of the standard sets

this limit:

^max <0-6

°ysi

ays2 .

Kq, (4)

where OyS1 and <7yS2 are the yield strengths at Tj and T2 and Kq is the

toughness at T 2

. Equation (5),

W “ °-6Kq, (5)

is a form of Eq. (4) where T x - T 2 and ays l /°ys 2

*- s 1* The intention is to

restrict K^x to a fraction of Kq, because if Kmav approaches Kj c too

closely, pronounced "pop -in" instabilities can affect the test records and

toughness measurements [30], increasing [31] or decreasing [32, 33] the

measured toughness. A decrease occurred for the affected test record shown

in Fig. 1(b).

For 13Mn steel, Eq. (4) implies a negligible effect of warm precracking

at 295 K on the toughness at 4 K if K^x < 14 MPa-m

1*. This closely agrees

with the experimental finding of Kjnax - 17 MPa-m^ (Fig. 4). These data

confirm that some limitation on K^x is necessary for cryogenic tests and

that Eq. (4) is sufficient to exclude WPS effects at 4 K when precracking is

conducted at 295 K. These are the first data ever put forth supporting Eq.

(4) for linear elastic tests near absolute zero.

Whereas Eq. (4) may be appropriate for Kj c tests, it is not useful for

ductile fracture tests, and it does not apply to the 22Mn steel. In fact,

the current standard for Jj c testing does not consider or recommended

guidelines for warm precracking of test specimens.

Limitations

The time consumed in precracking imposes a practical limitation. The

number of cycles required to generate a crack depends on notch type,

temperature, and technique (see section A2 of [14]), but the primary

253
variable is AK or Kmax . For steels, a realistic time period for crack

initiation and growth corresponds to 10 4 to 10 6 cycles. A typical Kmax to

accomplish this is 33 MPa ’in

15

.

Low Kmax may be favored to avoid load-history effects, but if Kmax is

too low, the crack will not initiate quickly. These two requirements are in

potential conflict, so it is impossible to properly precrack certain

materials at specific temperatures.

Extremely brittle materials are among those that may be impossible to

fatigue-crack in tension tests. Their Kjc values may be so low that Kmax

from Eq. (4) approaches or falls below the thresholds required for crack

initiation from the notch. This is common for ferritic steels tested at

cryogenic temperatures , and it may account for the problems encountered by

Stonesifer [34] in fatigue tests of A 533 B steel at 77 K. In some border

line cases, tensile precracking may be possible, but the time required is

prohibitive. For example, consider the 13Mn steel for precracking at 295

K, Eq. (4) requires a Kmax lower than 14 MPa-m^, but values below 14 MPa-m^

are too low to generate a crack in less than a million load cycles.

Problems can sometimes be resolved by precracking at an intermediate

temperature. For the example of 13Mn steel, precracking at 77 K may be

recommended. Precracking at 77 K is less conducive to a WPS effect at 4 K

than precracking at 295 K; the material's yield strength is much higher at

77 than at 295 K, and Kmax/ayS is also higher.

In summary, for Kj c tests at 4 K, precracking at 4 K is reliable but

costly. Precracking at 295 K is more convenient, but it will toughen some

steels. A third alternative is precracking at an intermediate temperature.

For tests of austenitic steels at 4 K, precracking at LNT works well because

liquid nitrogen is inexpensive and 4-K test specimens must be precooled in

liquid nitrogen anyway. For Jj c tests at cryogenic temperatures, there are

no guidelines yet to ensure freedom from the warm precracking effect;

therefore precracking at LNT or LHeT may be the best course to follow when

in doubt about the behavior of specific alloys.

CONCLUSIONS

1. Warm prestress effects do occur in some austenitic steels as well as

ferritic steels. In our study, a toughening effect from warm precracking occurred in the relatively brittle Fe-17Cr-3Ni-13Mn-0. 33N austenitic stainless steel at 4 K, but not in the higher toughness Fe13Cr-5Ni-22Mn-0 . 2 IN steel at 4 K.

2. Precracking at room temperature is difficult or impossible for Kjc tests of some low toughness steels at 4 K because of the warm prestress

effect. On the other hand, fatigue precracking at room temperature

appears to be acceptable for tough steels, such as 22Mn steel, that are

not susceptible to warm precracking effects when Jj c tests are performed at 4 K; more experience with other materials is needed.

3. Fatigue precracking at 77 K, using AK less than 36 MPa'm^5, appears to

be a good general procedure; so far it proved to be satisfactory for

254
five austenitic steels that have been subjected to Kj c or Jj c tests at

4 K. In this study, precracking at 77 K completely surpressed the warm

precracking effect in the susceptible 13Mn steel.

ACKNOWLEDGMENTS

This work was supported in the United States by the U. S. Department of

Energy, Office of Fusion Energy. D. Shepherd (NIST) provided fractographs

of the 13Mn steel.

REFERENCES

1. G.G. Chell, J.R. Haigh, and V. Vitek, A theory of warm pre-stressing,

experimental validation and the implications for elastic -plastic

failure criteria, C.E.R.L. Laboratory Note No. RD/L/N63/79 (1979).

2. D.A. Curry, A model for predicting the influence of warm prestress and

strain ageing on the cleavage fracture toughness of ferritic steels,

Int. J. Fract. . 22, 145-159 (1983).

3. R.W. Nichols, The use of overstressing techniques to reduce the risk of

subsequent brittle fracture, Parts 1 and 2, Brit. Weld. J

. Jan- Feb,

(1968)

4. B.W. Pickles, and A. Cowan, A review of warm prestressing studies, Int

J. Pres. Ves. & Pining . 14, 95-131 (1983).

5. A.J. Brothers and S. Yukawa, The effects of warm pre-stressing on notch

fracture strength, J . Basic Eng. . 85, 97-101, (1963).

6. T.C. Harrison and G.D. Feamehough, The influence of warm prestressing

on the brittle fracture cf structures containg sharp defects, J. Basic

Eng. . 94, 373-376 (1972).

7. R.W. Nichols, W.H. Levine, A. Quirk, and E.A. Bevitt, A limit approach

to the prevention of pressure vessel failure, Proc. First Int. Conf. on

Fracture, Sendai, Japan, 1673-1678 (1966).

8. H. Nakamura, H. Kobayashi, T. Kodaira, and H. Nakazawa, On the effects

of pre-loading on the fracture toughness of A 533B-1 steel, in:

Advances in Fracture Research. Proc. ICF 5 . 2, Pergamon, New York, 817-

824 (1981)

9. D.A. Curry, A micromechanistic approach to the warm prestressing of

ferritic steels, Int. J. Fract. . 17, 335-343 (1981).

10.

R.L. Tobler, Ductile fracture with serrations in AISI 310 S stainless

steel at liquid helium temperature, in: Elastic-Plastic Fracture:

Second Symposium. Vol. II — Fracture Resistance Curves and Engineering

Applications . ASTM STP 803, C.F. Shih and J.P. Gudas, eds., Amer. Soc.

Test. Maters., Philadelphia, 763-776 (1983).

255
11 . T. Ogata, K. Ishikawa, T. Yuri, R.L. Tobler, P.T. Purtscher, R.P. Reed,

T. Shoji, K. Nakano, and H. Takahashi, Effects of specimen size, sidegrooving, and precracking temperature on J- integral test results for

AISI 316 LN at 4 K, in: Adv . Crvo . Eng

.

. 34, A.F. Clark and R.P. Reed,

eds . , Plenum, New York, 259-266 (1988).

12. M. Shimada, R.L. Tobler, T. Shoji, and H. Takahashi, Size, sidegrooving, and fatigue precracking effects on the J- integral test

results for SUS 304 stainless steel at 4 K, in: Adv. Crvo. Eng. . 34,

A.F. Clark and R.P. Reed, eds., Plenum, New York, 251-258 (1988).

13. T. Horiuchi, R. Ogawa, and M. Shimada, Cryogenic Fe-Mn Austenitic

Steels, in: Adv. Crvo. Eng. . 32, R.P. Reed and A. F. Clark, eds.,

Plenum, New York, 33-42 (1986).

14. Standard test method for plane strain fracture toughness of metallic

materials, ASTM Designation: E 399-83, in: 1989 Annual Book of ASTM

Standards. Section 3. Metals Test Methods and Analytical Procedures . 03.01, Amer. Soc. Test. Maters., Philadelphia, 487-511 (1989).

15. Standard test method for Jjc» a measure of fracture toughness, ASTM

designation: E 813-81, in; 1986 Annual Book of ASTM Standards. Section

3. Metals Test Methods and Analytic Procedures . 03.01, Amer. Soc. Test.

Maters., Philadelphia, 968-990 (1986).

16. Standard test method for Jjc> a measure of fracture toughness, ASTM

designation: E 813-88, in: 1989 Annual Book of ASTM Standards. Section

3. Metals Test Methods and Analytical Procedures. 03.01, Amer. Soc.

Test. Maters., Philadelphia, 698-712 (1989).

17. Y.W. Cheng and D.T. Read, An automated fatigue crack growth rate test

system, in: Automated Test Methods for Fracture and Fatigue Crack

Growth . ASTM STP 877, W.H. Cullen, R.W. Landgraf, L.R. Kaisand, and

J.H. Underwood, eds., Amer. Soc. Test. Maters., Philadelphia, 213-223

(1985).

18. M. Shimada, R. Ogawa, T. Moriyama, and T. Horiuchi, Cryogenic Engineering . 21, 269-274 (1986), in Japanese.

19. R.L. Tobler and D.A. Meyn, Cleavage-like fracture in Fe-18Cr-3Ni-13Mn0.37N austenitic stainless steel at liquid helium temperature, Metall

.

Trans

.

. 19A, 1625-1631 (1988).

20. P.K. Liaw and J.D. Landes, Influence of prestrain history on fracture

toughness properties of steels, Metall. Trans. . 17A, 473-489 (1986).

21. J.L. Christian, J.D. Gruner, and L.D. Girton, Effects of cold rolling

on the mechanical properties of type 310 stainless steel at room and

cryogenic temperatures, Trans. ASM . 57, 199-207 (1964).

22. G.R. Irwin, Plastic zone near a crack and fracture toughness, in:

Mechanical and Metallurgical Behavior of Sheet Materials. Proc. 7th

Sagamore Ordnance Materials Research Conf

.

. U.S. Army Office of

Ordnance Research, 63-77 (1960).

256
23 . T.M. Banks and A. Garlick, The form of crack tip plastic zones, Eng

Fract. Mechs. . 19, 571-581 (1984).

24. C.M. Hsiao and E.J. Dulis, Trans ASM, 52 855-877 (1960).

25. J.C. Shyne, F.W. Schaller, and V.F. Zackay, Trans. ASM, 52 848-854

(1960).

26. J.D. Defilippi, K.G. Brickner, and E.M. Gilbert, Trans. AIME, 245 . 2141-2148 (1969).

27. G.R. Caskey, Jr., Report DP-MS-78-68, E.I. du Pont de Nemours and Co.,

Aiken SC, pp 1-21 (1978).

28. R.L. Tobler, R.P. Reed, and P.T. Purtscher, Linear elastic fracture of

high nitrogen austenitic stainless steels at liquid helium temperature,

J. Test. Eval. (1989).

29. T.A. Place, T.S. Sudarshan, C.K. Waters, and M.R. Louthan, Jr. Fractographic studies of the ductile- to-brittle transition in austenitic

stainless steels, in: Fractography of Modern Engineering Materials:

Composites and Metals, ASTM STP 948, Amer. Soc. Test. Maters.,

Philadelphia, 350-365 (1987).

30. W.F. Brown and J.E. Srawley, Plane Strain Crack Toughness Testing of

High Strength Metallic Materials . ASTM STP 460, Amer. Soc. Test.

Maters., Philadelphia, 1966, p. 51.

31. G. Clark, Significance of Fatigue Stress Intensity In Fracture Toughness Testing, Int. J. Fract. . 15, R179-R181 (1979).

32. T. Yokobori and T. Aizawa, Reports of the Research Institute for

Strength and Fracture of Materials . Tohoku University, 6, 19 (1970).

33. T. Kawasaki, T. Yokobori, Y. Sawaki, S. Nakanishi, and H. Izumi,

Fatigue Fracture Toughness and Fatigue Crack Propagation in 5.5% Ni

Steel At Low Temperature, in: Fracture 1977 . Vol. 3, D.M.R. Taplin,

Ed., University of Waterloo Press, Ontario, Canada, 1977.

34. F. Stonesifer, Effects of Grain Size and Temperature on Fatigue Crack

Propagation in A 533 B Steel, Eng. Fract. Mechs. . 10, 305-314 (1978).

257
'
JAPAN—U.S. COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

SIXTH DRAFT

OCTOBER 1988a

NOTE: This is a working document of a standard being developed by the

above committee for eventual submission to ASTM; it is intended

for committee review purposes, but not for general distribution

at this time.

Proposed Standard Method for

TENSILE TESTING OF STRUCTURAL ALLOYS AT LIQUID HELIUM TEMPERATURE

1 . Scope

1.1 This standard describes the procedures to be used for the tensile

testing of metals in liquid helium [4 or 4.2 K (—452°F)]. The format is

similar to that of other ASTM tensile test standards, but the contents include modifications for cryogenic testing, which requires special apparatus,

smaller specimen size, and concern for serrated yielding, adiabatic heating,

and strain- rate effects.

1.2 To conduct a tensile test by this standard, the specimen in a cryostat is fully submerged in normal liquid helium (He I) and tested using

crosshead displacement control at a moderate strain rate. Tests that use

load control or high strain rates are not considered in this standard.

1.3 This standard details methods for the measurement of yield strength,

tensile strength, elongation, and reduction of area. The determination of

the elastic modulus is treated in Method E 111.

1.4 Values stated in SI units are treated as primary; values stated in

U.S. customary units are treated as secondary.

1.5 This standard does not address the issue of safety associated with

cryogenic testing. It is the user's responsibility to learn and observe

appropriate safety practices. Safety guidelines for handling liquid helium

and other cryogens are available elsewhere (1).

NOTE 1—The boiling point of normal liquid helium (He I)

at sea level is 4.2 K. This value decreases slightly with

geographic elevation; it is 4.0 K at the National Institute of

Standards and Technology in Boulder, Colorado, 1677 m (5500

ft) above sea level. In this standard, the test temperature

is nominally designated 4 K.

259
2 . Applicable Documents

2 . 1 ASTM Standards

A 370 Methods and Definitions for Mechanical Testing of Steel Products

( Annual Book of ASTM Standards , Vol. 01.04)

E 4 Practices for Load Verification of Testing Machines

( Annual Book of ASTM Standards, Vols. 03.01, 04.02, 07.01, and

08.03)

E 6 Definition of Terms Relating to Methods of Mechanical Testing

( Annual Book of ASTM Standards

,

Vols. 03.01 and 08.03)

E 8 Methods for Tension Testing of Metallic Materials

( Annual Book of ASTM Standards, Vols. 01.02, 02.01, 02.02,

02.03, and 03.01)

E 8M Methods for Tension Testing of Metallic Materials, Metric

( Annual Book of ASTM Standards

,

Vol. 03.01)

E 29 Recommended Practice for Indicating Which Places of Figures

Are to Be Considered Significant in Specified Limiting Values

( Annual Book of ASTM Standards, Vols. 02.03, 03.01, 03.03,

03.05, and 14.02)

E 83 Practice for Verification and Classification of Extensometers

(Annual Book of ASTM Standards , Vol. 03.01)

E 111 Method for Young's Modulus, Tangent Modulus, and Chord Modulus

(Annual Book of ASTM Standards , Vol. 03.01)

2.2 Japanese Industrial Standards (JIS):

B 7721 Tensile Testing Machines

B 7728 Load Calibration Devices for Verifying Material Testing Machines

G 0567 Method of High Temperature Tensile Test for Steels and HeatResisting Alloys

Z 2201 Test Pieces for Tensile Test for Metallic Materials

Z 2241 Method of Tensile Test for Metallic Materials

Z 8103 Glossary of Terms Used in Instrumentation

Z 8401 Rules for Rounding Off of Numerical Values

NOTE 2—Japanese and English versions of these standards

are published by the Japanese Standards Association, 1-24,

Akasaka 4, Minato-ku, Tokyo, 107 Japan.

260
3. Definitions

3.1 The definitions of terms relating to tension testing that appear in

ASTM Standard E 6 shall apply here. The following definitions also apply:

3.1.1 tensile cryostat—a test apparatus for applying loads to test

specimens in cryogenic environments. A schematic illustration is shown in

Fig. 1.

3.1.2 Dewar—a vacuum- insulated container for cryogenic fluids.

3.1.3 adiabatic heating—the internal heating of a specimen resulting

from tensile testing under conditions such that the heat generated by plastic work cannot be quickly dissipated to the surrounding cryogen.

3.1.4 reduced section—section in the central portion of the specimen,

which has a cross section smaller than the gripped ends.

LOAD

I

Pul Rod

i j

Room Temperature

Load Frame

Alignment Device

Vacuum-insula

Transfer Tul

Cryoger

Load Frc

Spearr Vacuum-insulated

" Dewar

^

Electrical

Feed-Through

Vent

Dewar Seal

FIG 1 Example of a Cryostat with Double-Dewar Arrangement

for Tensile Testing at 4 K

261
3.1.5 length of the reduced section—-the distance between the tangent

points of the fillets that bound the reduced section.

3.1.6 adjusted length of the reduced section—the length of the reduced

section plus an amount calculated to compensate for strain in the fillet

region.

3.1.7 gage length—the original distance between gage marks made on the

specimen for determining elongation after fracture.

3.1.8 axial strain—the average of the strain measured on opposite

sides of the specimen at equal distances from the specimen axis.

3.1.9 bending strain—the difference between the strain at the surface

of the specimen and the axial strain. The bending strain varies from point

to point around and along the reduced section of the specimen.

3.1.10 maximum bending strain—the largest value of bending strain in

the reduced section of the specimen. It is calculated from strains measured

at three circumferential positions, at each of two different longitudinal

positions

.

3.1.11 discontinuous yielding stress, o £—the stress at the initiation

of the first measurable serration on the load-deflection curve. The

parameter a £ is a function of test variables and is not a material constant.

4. Significance and Use

4.1 In general, tensile tests provide information on the strength and

ductility of materials under uniaxial tensile stresses. This information

may be useful for alloy development, comparison and selection of materials,

and quality control. Under certain circumstances, the information may also

be useful for design.

4.2 At 4 K, the load—time and load—deflection record for metals tested

in displacement control are serrated (2). Serrations are formed by repeated

bursts of unstable plastic flow, followed by arrests. The unstable plastic

flow (or discontinuous yielding) is a free -running process that occurs

locally with sizable heat evolution at higher than nominal rates of strain.

Typical stress—strain curves for an austenitic stainless steel with serrated

yielding at 4 K are shown in Fig. 2.

4.3 Constant specimen temperature cannot be maintained at all times

during tests in a liquid helium environment. Owing to adiabatic heat, the

specimen temperature temporarily rises above 4 K during discontinuous

yielding events (see Fig. 2). This behavior varies with specimen size and

test speed, but altering the mechanical test variables cannot eliminate it

completely (3). Therefore, tensile property measurements of alloys in

liquid helium (especially ultimate strength, elongation, and reduction of

area) lack the full significance of property measurements at higher

temperatures where discontinuous yielding does not occur.

262
4.4 At 4 K, the stress—strain behavior of a material during unstable

plastic deformation depends on whether load control or displacement control

is used (4). Crosshead displacement control is specified because the goal

of this standard is material characterization by conventional methods. This

limitation must be taken into account when data are used for structural design in those applications where the actual circumstances may approach loadcontrolled conditions

* i i i r

:

e — 2.2 x 10- # s- 1

IMIIMM aaL/L100

tuc

D

50 £

UJ

CL2

4. 0.10 020 030

STRAIN

040

STRAIN STRAIN

FIG. 2 Typical Engineering Stress—Strain Curves and Specimen

Temperature Histories, at Four Different Nominal Strain

Rates, for AISI 304L Stainless Steel Tested in Liquid

Helium (Ref. 4)

263
5 . Apparatus

5.1 Test Machines—Machines used for tensile testing shall meet the

requirements of Practices E 4 regarding verification of load accuracy. The

machine compliance (displacement per unit load of the apparatus itself)

shall be measured. This can be done by removing the specimen from the load

train, or by replacing the specimen with a rigid block.

5.2 System Design—Alloy strengths often double or triple between room

temperature and 4 K. For the same specimen geometry, higher loads must be

applied to the cryostat, test specimen, load train members, and grips at

cryogenic temperatures. For most conventional test machines, which have a

maximum load capacity of 10 tons, it is recommended that the apparatus be

designed to accommodate one of the small specimens cited in section 7.2.2

of this standard.

5.3 Construction Materials—Many construction materials, including the

vast majority of ferritic steels, are brittle at 4 K. To resist embrittlement, the grips and other load- train members must be fabricated using

strong, tough, cryogenic alloys. Austenitic stainless steels (AISI 304LN)

,

maraging steels (200, 250, or 300 grades, with nickel plating to prevent

rust), and extra- low- interstitial (ELI) grade titanium alloys (Ti-6Al-4V

and Ti-5Al—2 . 5Sn) have been used, with proper design, grips, pull rods, and

cryostat frames.

5.4 Alignment:

5.4.1 Single- Specimen Apparatus—To avoid bending strains in the specimen, proper system alignment is essential. For a conventional singlespecimen test apparatus, the machine and grips shall be capable of loading

a precisely machined specimen so that the maximum bending strain does not

exceed 10% of the axial strain. This calculation is based on the strain

readings taken at zero load and at the highest load for which the machine

is being qualified.

NOTE 3—This requirement will minimize contributions from

the test apparatus to the bending strain. Tests performed

with a qualified apparatus may still vary in percent bending

strain owing to small variations in the specimens.

5.4.2 Multiple -Specimen Apparatus—For this type of cryostat, the

alignment depends on the type of fixtures used. The maximum bending strain

shall be measured and reported.

5.4.3 The testing apparatus may be qualified by axiality measurements

at room temperature and at 4 K. To perform axiality tests of the apparatus,

the specimen form should be the same as that used during cryogenic tests,

the specimen concentricity should be as nearly perfect as possible, and no

plastic strains should occur in the reduced section. In some cases, this

may necessitate the use of a relatively stiff, high-strength calibration

specimen.

264
5.4. 3.1 For round bar specimens, the maximum bending strain, as defined

in 3.1.10, is calculated from the strains measured at three circumferential

positions, at each of two different longitudinal positions. The strains are

measured with three electrical-resistance strain gages equally spaced around

the reduced section of the specimen. The two longitudinal positions should

be as far apart as possible, but not closer than one diameter to a fillet.

5. 4. 3.

For specimens of rectangular cross section, strain may be

measured at the center of each of the four sides, or in the case of thin

strips, near the outer edges of each of the two broad sides.

5. 4. 3.

3

To eliminate the effect of specimen bias, repeat the axiality

measurements with the specimen rotated 180 degrees, but with the grips and

pull rods retained in their original positions. The maximum bending strain

and strain at the specimen axis are then calculated as the average of the

two readings at the same position relative to the machine.

5.4.4 Strain Averaging Technique—Nonaxiality of loading is usually

sufficient to introduce errors in tensile tests at small strains when strain

is measured on only one side of the specimen. To rectify this, two extensometers or strain gages may be attached to opposite sides of the specimen.

The reported strain should be the average of the strains on both sides.

5.5 Gripping Mechanisms—The choice of gripping mechanism to be used

at 4 K is influenced by specimen type. Any suitable mechanism described in

Methods E 8 and E 8M may be used.

5.6 Dimension-Measuring Devices—Micrometers and devices used for measuring the dimensions of specimens shall be accurate and precise to at least

one-half the smallest unit to which a given dimension must be measured.

5.7 Cryostats and Support Apparatus

5.7.1 Cryostats—A cryostat capable of retaining liquid helium is prerequisite. In general, cryostat load frames for existing test machines must

be custom-built, but they may accommodate commercially available Dewars. The

cryostat may employ adjustable load columns to facilitate alignment. Several

practical designs, including turret-disc designs for multiple-specimen testing with a single cooling, are discussed in the literature (5-9).

5.7.2 Dewars—Stainless steel Dewars are recommended because they are

safer than glass Dewars and less expensive than fiberglass Dewars. A single

helium Dewar (see Fig. 1) is usually sufficient for short-term tests, but a double-Dewar arrangement, in which an outer Dewar of liquid nitrogen jackets

the inner Dewar of liquid helium, is also possible.

5.7.3 Ancillary Equipment—Dewars and transfer lines for liquid helium

must be vacuum insulated. Vacuum pumps, pressurized gas, and liquid nitrogen facilities are therefore required. After testing, the helium may be

released to the atmosphere, recycled as a gas, or reliquefied. Recycling

or reliquefaction requires large investments in purification and support

systems

265
5.8 Temperature Maintenance and Liquid-Level Indicators

5.8.1 Thermocouples—The intended temperature of 4 K is ensured by

maintaining a liquid helium environment. A thermocouple to measure the

specimen temperature is not required for routine tests.

5.8.2 Indicators—Although a thermocouple attached to the specimen is

not necessary, an indicator or meter is required to ensure that the specimen

remains fully submerged for the duration of testing. On—off indicators of

the carbon-resistor type may be used to verify that the liquid level always

exceeds some reference point above the specimen, or the liquid level may

be continuously monitored with superconducting wire sensors of appropriate

lengths positioned vertically inside the cryostat.

5.9 Strain Gages :

5.9.1 Selection—Strain-gage films bonded directly to the specimen

surface may be used to measure strain at 4 K (10). A satisfactory combination of gage active element, backing material, and bonding agent should

be selected on the basis of experience and manufacturer's recommendations.

One common choice is a Ni—Cr-alloy gage with a temperature -compensated

active element (7,11).

5.9.2 Characteristics—Strain gages are typically wired to a dummy

bridge using a three -wire temperature -compensating hookup, like that shown

in Fig. 3. The gage resistance is typically 120 or 350 0. A low excitation

voltage of about 1 to 2 V is recommended at 4 K to reduce Joule heating.

Typical full-scale operating ranges are 1% at room temperature and 2% at

4 K.

Input

Voltage

FIG 3 Method of Strain-Gage Wiring for Cryogenic Tests

266
5.9.3 Calibration—Strain-gage calibration at room temperature requires

calibration of the electronics to changes in resistance. Calibration at 4 K

requires consideration of the temperature dependence of the resistance in

the electronic calibration. Gage factors as a function of temperature are

provided by the manufacturers and in published research.

5.10 Clip-Gage Extensometers

5.10.1 Types—Detachable clip-gage extensometers for use at 4 K may be

built or purchased. An example is the beam gage, which uses four straingage films bonded in a Wheatstone bridge arrangement (11). Extension within

the specimen gage length is sensed by the extensometer , which is clipped to

retaining pins that are fixed to the specimen reduced section.

5.10.2 Characteristics—To measure the 0.2% offset yield strength,

a Class B—2 extensometer, as identified in Method E 83, may be used. The

extensometer shall meet the sensitivity and accuracy requirements of Method

E 83 and shall be tested to ensure accuracy at 4 K. Whenever possible, the

extensometer should be mounted directly to the specimen reduced section.

5.10.3 Calibration—Extensometers should be calibrated at room temperature and at 4 K with a suitable device. For calibrations at 4 K, a

micrometer with vertical extension tubes can be used with the extensometer(s) mounted at the lower end and immersed in liquid helium. Once the

calibration is known and proven to be accurate, linear, and reproducible,

room- temperature checks may be performed prior to each test series for indirect verification of the calibration at 4 K. Direct calibration at 4 K,

however, must be performed periodically, when damage is suspected or repairs

have been made

5.11 Capacitance Extensometers—Extensometers that use capacitance

measurement to monitor strain may be used (10). The type with overlapping

concentric cylinders has an extended strain range, an output that is linear

with displacement, and an adjustable sensitivity. The type with parallel

plates has high sensitivity, but the output must be compensated for the

hyperbolic dependence of the capacitance on displacement.

6. Sampling

6.1 Samples for tensile testing should be taken from the material in

its final condition to ensure that the properties measured are representative of the end product. Allowance should be made for any superficial

effects introduced by cutting operations.

6.2 Specimens should be cut from locations thought to be most representative of the stock material, realizing that test results for specimens

taken from selected locations of a part or material may not be representative of the whole. The conventional locations should normally be used:

6.2.1 For products 40 mm or less in thickness, diameter, or distance

between flats, the location should be at the center.

267
6.2.2

For products over 40 mm in thickness, diameter, or distance between flats, the location should be midway from the surface to the center.

6.3 The choice of specimen size and shape should be based on the requirements necessary to obtain representative samples of the material, and

on test-machine load capability (section 5.2).

6.4 The orientation of the specimen axis relative to the principal

working directions of the stock shall be specified using the notation in

section 4 of Method A 370.

7. Test Specimens

7 . 1 General

7.1.1 Types and Specifications—Any specimen configuration cited in

Methods E 8 or E 8M may be used. The specifications for dimensions,

tolerances, and surface finish ari stated in sections 6.1 through 6.17

of those standards.

7.1.2 Size—Specimens from sheet or wire products having relatively

small cross-sectional areas can be tested within the load limits of

conventional apparatus. Specimens from thick plate or bar products,

however, may need to be machined to a reduced cross-sectional area so

the load capacity of the machine is not exceeded.

7 . 2 Round Bar Specimens

7.2.1 Standard Room-Temperature Specimens—A 12.5-mm (0.5-in) diameter round bar specimen is the standard configuration for room- temperature

tests according to Methods E 8 and E 8M. Specimens of such strong alloys, however, require excessive loads at 4 K. For example, 210 kN is

required to test typical AISI 304LN steel at 4 K, whereas 100 kN or

10 tons is the limit for most machines.

7.2.2 Standard 4 -K Specimens—To meet the load limitations of conventional test machines, the round bar specimens in sections 7. 2. 2.1

and 7 .2.2.2 are recommended as standard for 4-K tests. The required

dimensions and tolerances for these specimens are given in Table 1.

Threaded or shouldered ends are common for gripping these specimens,

and the requirement of section 5.4.1 can be met by precise machining.

7.2.2.

1

Standard , small metric specimens . These specimens have a

7-mm diameter and a gage- length- to -diameter ratio of 5:1.

1 .2.2.2 Standard , small U.S. customary specimens

.

These specimens

have a 6.25-mm (0.25-in) diameter and a gage- length- to -diameter ratio

of 4:1.

268
7.2.3 Alternatives—If the 4-K standard specimens recommended above

are not appropriate, other sizes may be selected following the guidelines

of Methods E 8 and E 8M. The proportions of such specimens should be

similar to those of the standard specimens (see Fig. 4 of this standard

and Fig. 8 of Methods E 8 and E 8M)

TABLE 1 Standard Specimens for Room- Temperature Tests and

Recommended Proportionally Reduced, Standard

Small Specimens for 4-K Tests 1

(a)

Metric Versions Standard Specimen Standard

G/D ratio - 5 (room temperature) Small Specimen

(dimensions, mm) (4 K)

Nominal Diameter 12.5 7

G, gage length 62.5 ± 0.1 35 ± 0.1

D, diameter 12.5 ± 0.1 7 ± 0.1

R, fillet radius 10 7

A, reduced section 75 42

(b)

U.S. Customary Versions Standard Specimen Standard

G/D ratio - 4 (room temperature) Small Specimen

(dimensions, in) (4 K)

Nominal Diameter 0.5 0.25

G, gage length 2.000 ± 0.005 1.000 ± 0.005

D, diameter 0.500 ± 0.010 0.250 ± 0.005

R, fillet radius 0.375 0.1875

A, reduced section 2.25 1.25

x See also the notes to Fig. 8 in Methods E 8 and E 8M.

FIG 4 Round Bar Specimen Configuration (see Table 1)

269
7.2.4 Subsize specimens—Special care in fabrication and testing

is required for specimens with diameters less than 6 mm. As the specimen

size is reduced, factors such as machining, surface finish, alignment, and

the number of grains per cross section are of increasing importance.

8 . Procedures

8.1 Marking and Measuring the Test Specimen:

8.1.1 Gage Length—Gage length marks should be lightly punched,

scribed, or inked at appropriate locations on the reduced section of the

specimen, which is the conventional method. The gage length should normally

be five times the diameter for metric specimens or four times the diameter

for U.S. customary specimens. If another gage length is used, it must be

described in the report. Measure the gage length to the nearest 0.05 mm.

NOTE 4—For metals of low ductility, gage marks punched

or scribed on the reduced section may induce failure at those

locations due to stress concentrations. To avoid this, it may

be possible to coat the reduced section with layout ink, and

then mark the gage length by rotating the specimen in a jig

with knife edges scraping off the ink at the appropriate

intervals. Alternatively, gage marks may be placed on the

specimen shoulders, or the overall length of the specimen may

be measured.

8.1.2 Reduced Section—Measure the length of the reduced section and

the adjusted length of the reduced section, if applicable, to the nearest

0.05 mm.

8.1.3 Cross Section—Measure the cross-sectional area of the reduced

section as specified in section 7.1.1 of Methods E 8 and E 8M.

8.2 Specimen Installation—Install the specimen in the cryostat, leaving sufficient slack for instrumentation wires so they will not be stretched

or crimped during positioning of the Dewar and subsequent testing.

8.3 Seating and Alignment—If the gripping fixture involves loosefitting components, such as spherically seated bearings, prevent friction

or mismatch of the bearing surfaces at 4 K by first checking the seating

and alignment at room temperature. During this process of alignment, the

applied loads should be kept below one -third of the proportional limit of

the test material.

8.4 Cooling Procedure-—Remove any moisture from the apparatus prior

to cooling by drying it thoroughly with a pressurized air jet or heat gun.

Ice can block cryogenic transfer lines or cause erratic loading behavior if

it forms between various parts of the specimen, clip gage, and load train.

270
Next, position the Dewar and precool the apparatus to 77 K by transferring

liquid nitrogen into the cryostat. After attaining equilibrium at 77 K,

remove all the liquid nitrogen from the cryostat, and transfer liquid helium

into the cryostat until the specimen and grips are fully submerged. Testing

may begin after the system has reached thermal equilibrium at 4 K. The

specimen must remain fully submerged at all times during the test.

NOTE 5—The heat- transfer characteristics of gaseous

helium are inferior to those of liquid helium; therefore, a

liquid helium environment is required to minimize specimen

heating effects during discontinuous yielding.

8.5 Testing Speed:

8.5.1 Rate Control—Owing to adiabatic heating, tensile property

measurements at 4 K can be significantly affected by the testing speed.

Therefore, the test procedure must include a means of measuring and controlling the rate of crosshead motion. A nominal strain rate must be

specified, since the actual rate cannot be precisely controlled or maintained when discontinuous yielding occurs. The nominal strain rate is

calculated by dividing the crosshead rate by the length of the reduced

section. Alternatively, a pacing or indicating device may be used to

monitor the strain rate or an average strain rate may be determined by

observing the time required to effect a known increment of strain.

8.5.2 Rate Limit—The nominal strain rate at any time during the test

shall not exceed 1 x 10" 3 s

_1

. Higher rates may cause excessive specimen

temperature rises and therefore are not satisfactory for determining the

acceptability of materials.

8.5.3 Rate Change—The strain at which discontinuous yielding begins

usually increases with decreasing strain rate. If the first serration for

a given material occurs near 0.2% plastic strain, it may be desirable to

reduce the speed to prevent interference in the measurement of the yield

strength (see Fig. 5). Then, a relatively low strain rate may be used to

determine the yield strength, followed by a higher strain rate to complete

the test. Any convenient crosshead speed may be used up to a stress of onehalf the yield strength; after that, the crosshead speed must be such that

the nominal strain rate does not exceed the 1 x 10

3 s' 1 limit.

8.6 Measurement of Mechanical Properties

8.6.1 Load—Extension- Curve Method—To measure the yield strength, a

record of load versus extension must be obtained, up to at least 0.2%

plastic strain. The use of a strain measurement device for autographic

recording is recommended.

Measure the yield strength via the 0.2% offset method, following Methods

E 8 or E 8M, section 7.3.1. If the 0.2% offset line intersects the curve at

a load drop owing to discontinuous yielding, then report the highest stress

prior to that load drop as the yield strength.

271
A — Serrations after 0.2% strain B — Serrations before 0.2% strain

FIG 5 Stress-Strain Diagram for Determination of Yield Strength

by the Offset Method.

8.6.2 Load—Time -Curve Method—Yield strength measurements based on a

0.2% offset procedure applied to load-versus- time curves at 4 K are generally not recommended, but may be used for commercial test purposes with the

agreement of all parties involved. If this method is used, it should be

stated clearly in the report.

NOTE 6—Load—time curves for tests at 4 K are typically

nonlinear at the start and less regular than load—extension

curves. Also, the effective modulus of a thermally efficient

load train may be low and dependent on the liquid helium level

and the degree of temperature stabilization achieved throughout the system in the temperature range 295 to 4 K. As a result, yield strength data from load—time curves may be less

accurate than those of the recommended method.

8.6.3 Discontinuous yielding stress—Calculate the stress corresponding

to the point of initiation of discontinuous yielding by dividing the load at

the first measurable serration by the cross sectional area of the specimen.

8.6.4 Tensile Strength—Calculate the tensile strength by dividing the

maximum load carried by the specimen during the tensile test by the original

cross-sectional area of the specimen.

8.6.5 Elongation—Calculate the percentage increase of elongation

according to Methods E 8 or E 8M, section 7.6.

8.6.6 Reduction of Area—Calculate the percentage of reduction of area

according to Methods E 8 or E 8M, section 7.7.

8.6.7 Rounding Reported Test Data—Round-off the calculated numerical

test results according to Methods E 8 or E 8M, section 7.8.

272
8.6.8 Replacement Specimens—If necessary, discard any invalid data and

test replacement specimens as per Methods E 8 or E 8M, section 7.9.

9Report

9 . 1 General

9.1.1 Material Characterization—Describe the test material, including

manufacturing, processing, and metallurgical information.

9.1.2 Specimen Characterization—Describe the specimen location and its

orientation relative to the principal working directions of the stock. Also

report the specimen dimensions, including the cross-section dimensions, the

fillet radius, the reduced section length, and the adjusted length of the

reduced section (if used).

9.1.3 Strain Rate—Report the crosshead speed and the nominal strain

rate during yielding and after yielding, if a rate change is used.

9.1.4 Mechanical Property Measurements—Report the yield strength, the

ultimate strength, and the method of offset yield strength determination.

Include the method of extension measurement and the location of extensometer

attachment, if used. Also report the discontinuous yielding stress and the

strain rate at which it was measured, the tensile elongation and its method

of calculation, the gage -length- to -diameter ratio for round specimens, and

the reduction of area.

9.2 Optional Data—Report any optional data, such as measurements of

Young's modulus at 4 K, the mechanical properties at room temperature, and

the average grain size of the test material. Report the compliance of the

test machine with its cryostat.

9.3 Replicate Tests—If replicate specimens are tested, report the

number of tests, the average values of all mechanical property measurements,

and a measure of scatter.

9.4 Subsize Specimens—If subsize specimens are tested, state any

precautions taken with respect to specimen machining, surface condition, or

alignment, and report the grain size of the test material.

9.5 Anomalies—Report any anomalies in material behavior, test records,

or failure type and location.

10. Precision and Bias

10.1 Precision—The precision of these methods is being established by

a series of interlaboratory tests.

273
10.2 Bias—The bias of these methods includes quantitative estimates

of uncertainties of the dimension-measuring devices, the calibration of test

equipment, and the skill of the operators. At present, bias statements

should be limited to the documented performance of individual laboratories

.

REFERENCES

(1) Cryogenics Safety Manual—A Guide to Good Practice, 2nd Ed., Mechanical Engineering Publications, London, 1982.

(2) Basinski, Z. S., "The Instability of Plastic Flow of Metals at Very

Low Temperatures,” Proceedings of the Royal Society, Vol. A240, 1957,

pp. 229-242.

(3) Ogata, T. , Ishikawa, K. , and Nagai, K. , "Effects of Strain Rate on the

Tensile Behavior of Stainless Steels, Copper, and an Aluminum Alloy

at Cryogenic Temperatures," Tetsu-to-Hagane , Vol. 71, No. 10, 1985,

pp. 1390-1397.

(4) Ogata, T. and Ishikawa, K. "Time -dependent Deformation of Austenitic

Stainless Steels at Cryogenic Temperatures," Cryogenics, Vol. 26,

1986, pp. 365-369.

(5) Schwartzberg, F. R. , "Mechanical Property Testing Techniques for the

Cryogenic Temperature Range," in: Advances in Cryogenic Engineering,

Vol. 8, Plenum Press, New York, 1963, pp. 608—623.

(6) Evans, D. , Simmonds, G. E. , and Stapleton, G. B. , "Improved Facility

for Determining Mechanical Properties of Materials in Liquid Helium,"

in: Proceedings , Fourth International Cryogenic Engineering Conference, IPC Science and Technology Press, Guildford, Surrey, England,

1972, pp. 331-337.

(7) Hartwig, G. and Wuchner, F. , "Low Temperature Mechanical Testing

Machine," Review of Scientific Instruments , Vol. 46, 1975,

pp. 481—485.

(8) Novikov, N. V. "Mechanical Property Measurement Techniques of Structural Materials at Cryogenic Temperatures," in: Advances in Cryogenic

Engineering

,

Vol. 22, Plenum Press, New York, 1977, pp. 113-118.

(9) Horiuchi, T. , Shimada, M. , Fukutsuka, T. , and Tokuda, S., "Design and

Construction of an Apparatus for Testing Materials at Cryogenic

Temperatures," in: Proceedings, Fifth International Cryogenic

Engineering Conference

,

IPC Science and Technology Press, Guildford,

Surrey, England, 1977, pp. 465—468.

(10)

Sparks, L. L. , "Temperature, Strain, and Magnetic Field Measurements,"

in: Materials at Low Temperatures, R. P. Reed and A. F. Clark, eds .

,

American Society for Metals, Metals Park, Ohio, 1983.

274
(11) Read, D. T. and Tobler, R. L. , "Mechanical Property Measurements at

Low Temperatures," in: Advances in Cryogenic Engineering—Materials,

Vol. 28, Plenum Press, New York, 1982, pp. 17-28.
wm
TECHNOLOGY TRANSFER

277

TECHNOLOGY TRANSFER PROGRAM

LEADER: N.J. Simon

CONTRIBUTORS: E.S. Drexler, R.P. Reed

OBJECTIVES

o Organization of workshops to promote interaction between designers and

material specialists, to discuss issues related to low- temperature

material needs of the fusion energy program, and to present new lowtemperature data for structural alloys, composites, and weldments.

o Preparation of an annual report, as well as monthly highlight reports

to the Office of Fusion Energy, U.S. Department of Energy.

o Evaluation of low- temperature mechanical and physical properties of

insulator, conductor, and structural materials for cryogenic copper and

superconducting magnets; preparation of handbook pages and supporting

documentation; distribution of handbook pages to participants in the

fusion energy program and to the Materials Handbook for Fusion Energy

Systems

RESEARCH HIGHLIGHTS

o The Eleventh Annual Cryogenic Structural Materials Workshop was

organized; it was held October 8-9, 1988 in Colorado Springs, Colorado.

o "Materials Studies for Magnetic Fusion Energy Applications - XI",

(NBSIR 88-3082, 398 pages, 1988, editor R. P. Reed) was prepared,

published, and distributed.

o Data were collected for handbook pages on C10100—C10700 copper and

C17000-C17510 beryllium copper. Over one thousand documents (reports

and journal articles) were acquired, coded for property information,

and entered into a data-base management system. Evaluation of

documents pertaining to C10100-C10700 copper tensile, elastic,

electromagnetic, fatigue, and thermal properties was completed; data

were extracted and analyzed; handbook pages on these properties were

prepared and distributed. Evaluation of documents for similar

properties of C17000-C17510 beryllium copper was also completed; data

were extracted and analyzed; handbook pages on tensile, elastic,

electromagnetic, fatigue, and thermal properties were prepared and

distributed.

o Handbook pages on C50500—C52400 phosphor bronze tensile, elastic,

fatigue, electro-magnetic, and thermal properties are in preparation.

Documents (over 250) have been acquired, and preliminary data sets

extracted and analyzed.

279

ELEVENTH CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

Clarion Hotel

Colorado Springs, Colorado

October 18-19, 1988

This workshop provided an opportunity for discussion of problems,

advances, and goals in the development of low- temperature

materials for cryogenic and superconducting magnets for fusion

energy systems. The 1988 workshop emphasized:

• Cryogenic properties of copper alloys and laminates

pertaining to design requirements for the compact ignition

tokamak (CIT)

• Identification of materials needs for future magnet systems

for the International Thermonuclear Experimental Reactor

(ITER)

• New low- temperature materials developments in the United

States and international research cooperation

281
11TH CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

FRQGRAM

TUESDAY. OCTOBER 18

jl

8:00 a.m. Continental Breakfast

8

8:30 a .m. Introduction and Announcements R. Reed, NIST

S GENERAL 8 8:40 a .

m

. DOE Overview and Perspectives A. Opdenaker, OFE

9:00 a.m. CIT Design Overview and Materials Program C. Bushnell, PPPL

9:30 a.m. Cryogenic Materials for the International

Thermonuclear Experimental Reactor (ITER)

L. Summers, J. Mill

LLNL

WELDING

10:00 a.m. Cryogenic Material Properties of Stainless

Steel Tube- to -Flange Welds

T. Siewert, C. McCo't

D. Vigliotti , NIST

10:25 a.m. Joining of Nonferrous Materials G. Goodwin, ORNL

10:50 a.m. Role of Inclusions in the Fracture of

Austenitic Welds

C. McCowan, T. Siewit

NIST

11:15 a.m. FN Predictive Diagram for Stainless Steel

Welds

T. Siewert, C. McCoi:

NIST, D. Olson, CSM

•—

Lunch Break

INSULATORS

1:00 p .m. SSC Dipole Magnet Design and Performance Update T. Nicol, Fermilab

1:30 p . m

.

Cryogenic Torsional Shear and Fracture Strength

of Epoxy and Polyimide Composites

M. Kasen, N. Munshi,

CTD

1:50 p.m. Radiation Effects in Structural Insulators R. Reed, NIST

2:30 p.m. Discussion of U.S. /Japan Efforts to Develop

Radiation-Resistant Insulators

M. Kasen, CTD

2:50 p.m. Insulation Development for ITER N. Munshi, M. Kasen,

3:10 p.m. CIT Insulation R & D Program T. McManamy, MartinMarietta, ORNL

3:45 p.m. Break

4:00 p.m. Discussion of CIT Program

6:00-7:30 Reception

282
WEDNESDAY. OCTOBER 19

8:00 a.m. Continental Breakfast

STRUCTURAL

8:30 a.m.

9:10 a.m.

9:40 a.m.

10:10 a.m.

10:40 a.m.

11:10 a.m.

11:40 a.m.

1 : 30 p .m. 2:00 p.m.

2:30 p.m.

3:00 p.m.

3:30 p.m.

4:00 p.m.

4:30 p.m.

ALLOYS

316LN Forgings for C-MOD-Metallurgical and

Design Considerations

Metallurgical Factors that Influence Fracture

Toughness of Austenitic Steels

Fracture Behavior of 304 Stainless in an 8 T

Field

Creep of Copper and Austenitic Steels; 4 to 295 K

Effect of Nitrogen on the LCF Properties of

316 Stainless Steel at 300 K and 77 K

Load- Control Tensile Measurements on Austenitic

Steels

Discussion of Load- Control Tensile Measurements

in Terms of Magnet Design

Lunch Break

Review of International Cooperative Programs

U.S. -Japan Program for Mechanical Property Test

Standards - An Update

Fatigue Crack Growth in Metastable Austenitic

Stainless Steels

The Bureau of Mines Development Program for New

Austenitic Stainless Steels and Cryogenic

Characterization

The Cryogenic Mechanical Properties of Advanced

Aluminum Alloys

Fatigue Properties of High Strength Aluminum

Alloys

Handbook - Structural Materials for

Superconducting Magnets - Progress Report

R. Ballinger, H. Becker,

MIT

P . Purtscher , NIST

J. Chan, J. Glazer, and

J. Morris, Jr., U. CA

R. Walsh, N. Simon,

C. King, R. Reed, NIST,

J. Arvidson, MRE, Inc.

J.-B. Vogt, J. Foct,

U. Lille

H. Lee, J. Han,

R. Reed, NIST

R. Reed, NIST

R. Reed, NIST

R. Tobler , NIST

Z. Mei, J. Morris, Jr.,

U. CA

J. Dunning, M. Glenn,

USBM

P. Purtscher, R. Walsh,

T. Siewert, R. Reed, NIST

J . Morris , Jr .

,

J. Glazer, U.CA

J. Han, L. Ma,

R. Tobler, NIST

N. Simon, L. Delgado,

C. King, R. Reed, NIST

283
ELEVENTH CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

SUMMARY OF TECHNICAL PRESENTATIONS

GENERAL SESSION

A. Opdenaker from the DOE Office of Fusion Energy presented an

overview of DOE perspectives on fusion. First, he outlined the

overall budget for magnetically contained fusion energy systems. Over

half of this budget goes to confinement system design, about one

fourth to applied plasma physics and about one sixth to development

and technology. The latter category includes the materials research

for superconducting and cryogenic magnets that was the subject of this

workshop. The current prototype reactor is the Compact Ignition

Tokamak (CIT) , which will have normal copper or copper alloy coils

operating at 77 K to produce fields up to 16 Tesla. The next

prototype, the International Thermonuclear Experimental Reactor

(ITER) , will have superconducting coils that operate at 4 K with

fields of 11 — 12 Tesla. ARIES, a future design, may have fields as

large as 24 — 30 Tesla. Priorities and future direction of the DOE

fusion energy program will be 1) to understand confinement physics,

since there is still a degree of skepticism due to uncertainty in

confinement mechanisms; 2) base a practical research program chiefly

on the Tokamak design; 3) since budgets will be level at best, old

machines will have to be given up in favor of new ones and the next

series of machines will have to be built with international

cooperation and research assistance from universities; 4) safety and

environmental concerns must be addressed.

C. Bushnell of PPPL gave an overview of the CIT design and

materials program. There will be about 420 toroidal coil plates of

Inconel 718 laminates, C17510 beryllium copper, or C15725 aluminumdispersion strengthened copper. The latter material is also an

alternative for the poloidal coils. The toroidal field pulse will

have a 12 s rise, a constant ohmic heating period of 5 s , and an 8 -

12 s decay. Part of the system is expected to warm up nearly to room

temperature during a pulse. The device is to be designed for a total

of about 3000 pulses. A cooldown period of about one h is expected.

Techniques for explosive bonding of copper to Inconel have been

investigated. Coil development will take place through 1991, and the

device is expected to be operational with plasma ignition in 1996.

The U.S., USSR, Japan and the European Community are involved in

a three-year effort to design ITER. Materials design issues for the

reactor were described by L. Summers of LLNL. ITER will be a very

large machine, with each of 16 toroidal field coils weighing 350

tonnes, and a total cold mass of 10,000 tonnes of superconductor,

sheathing and magnet case. A cable- in-conduit configuration with an

Nb 3 Sn superconductor will be contained in a case of 304LN or 316LN

stainless steel. The magnet will be pulsed, and therefore fatigue

284
problems are a serious materials issue. However, the magnet

insulation will receive a total radiation dose of only about 107 Gy,

which is 100 times smaller than the radiation dose predicted for

earlier, more compact reactor designs such as TIBER II. Since vacuum

impregnation of the coils with the insulating material is a

requirement, an epoxy matrix rather than a polyimide or thermoplastic

matrix is likely to be chosen. The problem of scrubbing of this

insulation as the coils are pulsed will have to be investigated.

Another important materials issue is the development of fabrication

methods, including welding and inspection, for the conduit and

superconductor. Conduit materials under consideration are Inconel

908, JBK 75, and A- 286. Selection of this alloy is restricted by the

requirement that it retain good properties after the Nb 3 Sn heat

treatment

WELDING SESSION

The three reports given in the welding session were presented by

NIST staff (T.A. Siewert and C.N. MeCowan ) . They reported recent

advances in predicting the ferritic phase content, solidification

mode, fatigue life, and fracture toughness of stainless steel welds.

The data generated by these programs are expected to be useful to

fusion magnet designers. Data on increasing weld metal toughness by

reducing the inclusion content of welds were of particular interest.

From a practical standpoint, production welding processes must be used

during the fabrication of fusion magnet structures and the NIST

studies continue to address means to optimize the toughness of welds

using modified production techniques.

INSULATOR SESSION

The session on structural composites, following the trends of all

recent workshops, again emphasized the concern for magnet insulation

in a radiation environment at low temperatures and the scarcity of

appropriate data. The proposed closing of the Oak Ridge LTNIF

facilities for the Bulk Shielding Rector has completely frustrated the

ongoing programs scheduled to use the ORNL facilities to develop

radiation- resistant structural insulators. Past data on this subject

were reviewed by R. Reed (NIST)

The CIT program to characterize structural insulators for the

77 - 300 K environment with a pulsed magnet was extensively discussed

(T. McManamy, ORNL) Current plans for composite supports for the SCC

program were presented by T. Nicol (Fermilab).

Generally, there are two types of composites using either 2- or

3- dimensional weaves of S-2 glass with either epoxy, bismaleimide , or

polyimide resins. Since the cost varies by at least an order of

magnitude, depending on the resin and reinforcement system, and since

little or no useful data have been generated that simulates operating

conditions, extensive research is required on this subject.

285
STRUCTURAL ALLOYS - SESSION I

In the first session on structural alloys, mechanical property

and metallurgical research from MIT, LBL, University of Lille in

France, as well as from NIST was summarized. The talks demonstrated

the diverse metallurgical/mechanical factors that affect material

performance at cryogenic temperatures. There are still technical

questions to answer before we can confidently build a safe magnet case

to operate in liquid helium. Lee (NIST) described how loading rate

affects the flow stress of different austenitic steels. The ultimate

stress of the steels was reduced when the loading rate was increased

to values that were similar to the loading rate a magnet might

experience. This behavior may change the limiting design stresses for

magnets. Chan from LBL described the results of their fracture

toughness tests in the presence of high magnetic fields. No definite

conclusions can be drawn from LBL's results at this point. The

relatively low values of fracture toughness of 304 steel at 4 K would

appear to be more significant than any effect due to the magnetic

field. The global nature of the research effort in structural alloys

was evident at this session. The large 316LN forgings for the C-MOD

were made in West Germany because no company in the U.S.A. had the

capability. Professor Foct from France showed some interesting

effects of nitrogen additions on the deformation behavior of 316- type

steels

.

STRUCTURAL ALLOYS - SESSION II

This session covered standards, alloy, and handbook developments.

R. Reed, R. Tobler and N. Simon (NIST) described the handbook and

standards related activities in the past year. Rapid progress has

been made in the development of tensile and fracture toughness test

standards because of the VAMAS and U.S. — Japan international

programs. The official processing of draft standards through ASTM

should be completed by 1990. Regarding steel developments, J. Dunning

(U.S. Bureau of Mines) and P. Purtscher (NIST) described a cooperative

program between the USBM and NIST to produce good, low cost austenitic

stainless steels containing only 9%Cr. One of the four candidate

materials (Fe-9Cr-14Ni-4Mn-2Mo) has 4-K yield strength and fracture

toughness properties equivalent to commercial 304 steels, proving that

the new low Cr alloys may be practical. Similarly, J.W. Morris (LBL)

and J. Han (NIST) considered some new aluminum alloys. Previous data

demonstrated that the Al-Li alloy 2090 in plate form surpasses

conventional 2000 series alloys in strength- toughness properties at

cryogenic temperatures. Now it appears that Al-Sc alloys surpass the

Al-Li alloys in thin sheet tearing resistance at cryogenic

temperatures. However, other data were presented, indicating that the

superiority of 2090 versus alloys such as 2219 does not extend to thin

sheet materials tested in notched tension or short term fatigue tests

at cryogenic temperatures

.

286
N8T-I14A U.S. DEPARTMENT OF COMMERCE i. now ON IVONT 1 REV. S-BB) NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY

BIBLIOGRAPHIC DATA SHEET

NISTIR 89-3931

2. PERPORMMQ OROAMZATKM REPORT NUMBER

X PUBLICATION DATE

November 1989

TTTLE AMO SUBTITLE

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures — XII

AUTHOR(3)

R.P. Reed and R.L. Tobler, Editors

PERFORMING ORGANIZATION (IF JOINT OR OTHER THAN NIST, SEE INSTRUCTIONS)

U.S. DEPARTMENT OF COMMERCE

NATIONAL INSTITUTE OF STANDARDS ANO TECHNOLOGY

GAITHERSBURG, MO 20899

7. CONTRACT/GRANT NUMBER

X TYPE OF REPORT ANO PERIOD COVERED

SPONSORING organization name ANO COMPLETE ADORESS (STREET, CITY, STATE. ZIP)

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

0. SUPPLEMENTARY NOTES

Previous reports in this series: NBSIR 88-3082, NBSIR 87-3067, NBSIR 86-3050,

NBSIR 85-3025, NBSIR 84-3000, NBSIR 83-1690, NBSIR 82-1667, NBSIR 81-1645, NBSIR 80-1627,

NBSIR 79-1609, NBSIR 78-884

DOCUMENT DESCRIBES A COMPUTE* PROGRAM; SP-1 85, PIPS SOFTWARE SUMMARY, IS ATTACHED.

II. ABSTRACT (A 200-WORD OR LESS FACTUAL SUMMARY OF MOST SIGNIFICANT INFORMATION. IF DOCUMENT MCLUOES A SIGNIFICANT BIBLIOGRAPHY OR

LITERATURE SURVEY, MENTION IT HERE.)

This report contains the results of a research program to determine the properties of

materials that may be used in cryogenic structures for the superconducting magnets of

fusion energy power plants and prototypes. Its purpose is to facilitate their design

and development. The program was developed jointly by the staffs of the National

Institute of Standards and Technology and the Office of Fusion Energy of the Department

of Energy; it is managed by NIST and sponsored by DOE. Research is conducted at NIST

and at other laboratories through subcontracts with NIST.

Research results for 1988 are presented in technical papers under four headings that

reflect the main program areas: Structural Alloys, Welding, Technology Transfer, and

United States-Japan Development of Test Standards. Objectives and research highlights

are summarized in the introduction to each program area.

12. KEY WORDS (« TO 12 ENTRIES; ALPHABETICAL ORDER; CAPITALIZE ONLY PROPER NAMES; ANO SEPARATE KEY WORDS EY SEMICOLONS)

austenitic steels; composites; cryogenic properties; elastic properties; low temperatures;

mechanical properties; nonme tallies; stainless steels; structural alloys; welds

13. AVAILABILITY 14. NUMBER OF PRINTED PAGES

XX UNLIMITED 277

FOR OFFICIAL DISTRIBUTION. DO NOT RELEASE TO NATIONAL TECHNICAL BIFORMATION SERVICE (NT1S).

ORDER FROM SUPERINTENDENT OF DOCUMENTS, U.S. GOVERNMENT PRINTING OFFICE,

WASHINGTON. DC 20402.

IS. PRICE

A13

XX ORDER FROM NATIONAL TECHNICAL INFORMATION SERVICE (NTIS). SPRINGFIELD, VA 22181. ELECTRONIC FORM
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