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NBSIR 88-3082

TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

applications at low temperatures -XI

Department of Energy

I Office of Fusion Energy

Washington, D.C. 20545 Prepared by

Fracture and Deformation Divisioi

National Bureau of Standards

Boulder, Colorado 80303

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MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

applications at low TEMPERATURES -XI

Edited by

R P Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

May 1988

Sponsored by

Department of Energy

Office of Fusion Energy

Stimulating America's Progress 1913-1988

U.S. DEPARTMENT OF COMMERCE, C. William Verity, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director
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COHTEHTS

SUMMARY vii

ORGANIZATIONAL CONTACTS ..... viii

PROGRAM DESCRIPTION ix

STRUCTURAL ALLOYS

Structural Alloys Program 3 Summary of Mechanical Property Measurements of Structural Alloys ... 5 Austenitic Stainless Steels with Emphasis on Strength

at Low Temperatures 9 R. P. Reed

Low- Temperature Properties of High -Manganese Austenitic Steels .... 47

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

Design of 316LN-Type Alloys 71

N. J. Simon and R. P. Reed

Cleavage- like Fracture Along Slip Planes in Fe—18Cr—3Ni—13Mn—0 . 37N

Austenitic Stainless Steel at Liquid-Helium Temperature 89

R. L. Tobler and D. Meyn

Fracture Behavior of 316LN Alloy in Uniaxial Tension

at Cryogenic Temperatures 103

P. T. Purtscher, R. P. Walsh, and R. P. Reed

Effect of Void Nucleation on Fracture Toughness

of High-Strength Austenitic Steels 113

P. T. Purtscher, R. P. Reed, and D. T. Read

Low- Temperature Phase and Magnetic Interactions

in Face-Centered-Cubic Fe—Cr—Ni Alloys 131

C. Almasan, T. Datta, R. D. Edge, E. R. Jones, J. W. Cable,

and H. Ledbetter

Effects of Grain Size and Cold Rolling on Cryogenic Properties of Copper 147

R. P. Reed, R. P. Walsh, and F. R. Fickett

Tensile and Creep—Fatigue Properties of

a Copper—Stainless-Steel Laminate 159

L. L. Scull and R. P. Reed

iii
WELDING

Welding Program 171

Summary of Mechanical Property Measurements of Welds 173

Ferrite Number Prediction to 100 FN in Stainless Steel Welds .... 175

T. A. Siewert, C. N. McCowan, and D. L. Olson

Mechanical Properties of Electron Beam Welds in Thick Copper .... 201

T. A. Siewert and D. P. Vigliotti

Cryogenic Material Properties of Stainless Steel

Tube- to -Flange Welds . ........... 215

T. A. Siewert, C. N. McCowan, D. P. Vigliotti

Strength and Toughness of an Aluminum- Deoxidized Weld

for Cryogenic Service ....... .... 237

C. N. McCowan, T. A. Siewert, and D. J. Kotecki

Inclusions and Fracture Touthness in Stainless Steel Welds at 4 K . . 247

C. N. McCowan and T. A. Siewert

NONMETALLIC

S

Nonmetallics Program ........................ 259

Internal-Friction Spectrum of a Glass—Epoxy Composite ....... 261

M. Weller and H. Ledbetter

Torsion Modulus and Internal Friction of

a Fiber-Reinforced Composite ... ..... 269

H. Ledbetter and Ming Lei

TECHNOLOGY TRANSFER

Technology Transfer Program ..... 277

Tenth Annual Cryogenic Structural Materials Workshop ........ 279

N . J . S imon

UNITED STATES-JAPAN DEVELOPMENT OF TEST STANDARDS

United States—Japan Development of Test Standards Program ...... 289

Trip Report — ASTM Technical Committee Meetings . 291

R. L. Tobler

Loading-Rate Effects on Discontinuous Deformation

in Load-Control Tensile Tests ..... ....... 295

T. Ogata, K. Ishikawa, R. P. Reed, and R. P. Walsh

iv
Round-Robin Tensile and Fracture Test Results for

an Fe—22Mn—13Cr—5Ni Austenitic Stainless Steel at 4 K 305

H. Nakajima, K. Yoshida, S. Shimamoto, R. L. Tobler,

P. T. Purtscher, and R. P. Reed

Size, Side-Grooving, and Fatigue -Precracking Effects

on J- integral Test Results for SUS 304 Stainless Steel at 4 K . . . . 317

M. Shimada, R. L. Tobler, T. Shoji, and H. Takahashi

Effects of Specimen Size, Side Grooving, and Precracking

Temperature on J- integral Test Results for AISI 316LN at 4 K . . . . 329

T. Ogata, K. Ishikawa, T. Yuri, R. L. Tobler, P. T. Purtscher,

R. P. Reed, T. Shoji, K. Nakano, and H. Takahashi

Mechanical Tests of Large Specimens at 4 K:

Facilities and Results 341

K. Yoshida, H. Nakajima, M. Oshikiri, R. L. Tobler,

S. Shimamoto, R. Miura, and J. Ishizaka

Acoustic Emission and Its Applications to Fracture Studies

of Austenitic Stainless Steels at 4 K 351

H. Takahashi, T. Shoji, and R. L. Tobler

Proposed Standard Methods for Fracture Toughness Testing

of Austenitic Stainless Steels at Liquid Helium Temperature

(fifth draft) 363

R. L. Tobler

Proposed Standard Method for Tensile Testing

of Structural Alloys at Liquid Helium Temperature (fifth draft) . . . 383

R. L. Tobler

v

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 magnetic

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 Bureau of Standards and the

Office of Fusion Energy of the Department of Energy; it is managed

by NBS and sponsored by DOE. Research is conducted at NBS and at

other laboratories through subcontracts with NBS.

Research results for 1987 are presented in technical papers

under five headings that reflect the main program areas: Structural Alloys, Welding, Nonme tallies , 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

Bureau of Standards, nor does it imply that the material or equipment identified is necessarily the best available for the purpose.

vii
ORGANIZATIONAL CONTACTS

Specific technical questions may be directed to the following

people, who contributed to major aspects of the program during the

fiscal year 1987.

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

Program Monitor A. L. Opdenaker (301) 353-4954

National Bureau of Standards, Boulder, Colorado 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. McGowan (303) 497-3699

T. A. Siewert (303) 497-3523

Nonmetallics H. Ledbetter (303) 497-3443

Technology Transfer R. P. Reed (303) 497-3870

N. J. Simon (303) 497-3687

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

Development of R. P. Reed (303) 497-3870

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

R. P. Walsh (303) 497-3818

viii
PROGRAM DESCRIPTION

The overall 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 specific 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, con- struction, and measurement programs, such as the Large Coil Project,

MFTF-B, and ITER (International Thermonuclear Experimental Reactor),

we are aware of new problems as they arise . This year ' s report contains results of continued research on copper and copper alloys, 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.

ix

STRUCTURAL ALLOYS

1

STRUCTURAL ALLOyS PROGRAM

LEADER:

CONTRIBUTORS

R. P. Reed

M. W. Austin, L. A. Delgado, H. Ledbetter, P. T. Purtscher,

D. T. Read, N. J. Simon, R. P. Walsh

OBJECTIVES

• Development of strong, tough structural alloys for use in superconducting

magnets

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

pulsed superconducting magnets

• Development of effective codes and standards for low- temperature property

measurements and structural design

RESEARCH HIGHLIGHTS

• Two new reviews have been written:

Austenitic Stainless Steels with Emphasis on Their Low -Temperature

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

• 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,

316LN, and 316LHN 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 316LN offers higher strength than 304LN with equivalent toughness. Nitrogen additions permit the attainment of high yieldstrength levels with little added cost.

3
• 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 K Elastic

295-4 K

Fracture Toughness

295, 76, 4 X

Fatigue Crack

Growth Rate

295, 76, 4 K

Fatigue

S-N

4 K

Report^

Volume

(page)

AUSTENITIC STEELS

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

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

Fe-13Cr-19Mn USSR III VI III (79)

VI (41)

Fe-l3Cr-22Mn-0.2N Japan VI VII (65)

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

Fe-17Cr-18Mn-0.5N U.S. VII, VIII VII (65) VIII VIII (181)

Fe-17Cr-9Ni-8Mn U.S. III III (91)

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

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

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

IX, X rv (147) V (185)

VI (157) — VIII (181,209,251)

IX (15,27,43)

X (77)

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

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

Fe-18Cr-(ll-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 I, VI, VII I (93) VIII III (91)

VI (53) VII (85) VIII (167,219)

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

(0-6)Mn

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

4Mn-(0 . 1-0 . 3 )N IX (53)

Fe-19Cr-9Ni U.S. I, II, I, II. 1,11, II, III, IV I (15,71,213)

[304] m,rv, Ill, IV, III, IV, II (79,149,175)

vi, VII, VI, VII IX III (15. 91. 105, n;

IX, X X rv (37,101.203,

215,227)

V (71)

vi (73) VII (13.47,75.157

IX (27)

X (77)

*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).

5
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

Alloy

[Designation]

Country

of

Supplier

Tensile

295-4 K Elastic295-4 K

Fracture Toughness

295, 76, 4 K

Fatigue Crack

Growth Rate

295, 76, 4 K

Fatigue

S-N

4 K

Report^

Volume

(page)

Fe-19Cr-9Ni O.S. II, III, 1,11,1V II, IV III. IV, I (213)

[304L] rv.vin, VIII, IX VIII II (79,123,175)

IX, X IV (37,101,131,

215,302)

VIII (181,229,251)

IX (27,141)

X (9)

Fe-19Cr-9Ni- u.s. IV, V, V IV, V, VIII, rv (77)

(l-10)Mn VIII IX V (15,59,189)

VIII (181)

IX (53)

Fe-19Cr-9Ni-N O.S. I, III, IV I, III, IV, I, III, IV i (93)

[304N] IV, IX IX hi (15) IV (37,101,203,

215) IX (27)

Fe-19Cr-9Nl-N O.S. II, III, IV, VIII II, III, IV, II, III, IV, ii (35,79)

[ 304LN] rv,v,vi VIII, IX VIII hi (15)

VIII, IX IV (37,101,203)

V (29) VI (113)

-- VIII (181,207,229)

IX (159)

Fe-19Cr-10Ni- O.S. III, VIII VIII, IX III, VIII III III (15)

(0.03-0.29)N- VIII (145,181)

(0 .03-0 .09 )C IX (159)

Fe-20Cr-7Ni-8Mn O.S. VIII VIII VIII (151,181)

[216]

Fe-20Cr-8Ni-2Mn O.S. VIII VIII VIII (151,181)

[304HN]

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

VI (113)

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

III (91)

VI (157) VIII (151,181)

Fe-21Cr-l2Ni-5Mn O.S. i , ii III 1,11 1,11 I (93)

II (79) III (91)

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

[310] VII, V (71,145)

VIII, VI (73)

IX, X VII (65) VIII (181)

IX (95) X (9)

^Materials Studies for Magnetic Fusion Energy Applications at NBSIR 78-884 (I); NBSIR 79-1609 (II); NBSIR 80-1627 (III);

NBSIR 82-1667 (V); NBSIR 83-1690 (VI); NBSIR 84-3000 (VII);

NBSIR 86-3050 (IX); NBSIR 87-3067 (X).

Lour Temperatures

NBSIR 81-1645 (IV);

NBSIR 85-3025 (VIII)

;

6
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

Country Fracture Fatigue Crack Fatigue Report^

Alloy of Tensile Elastic Toughness Growth Rate S-N Volume

[Designation] Supplier 295-4 K 295-4 K 295, 76, 4 K 295, 76, 4 K 295, 4 K (page)

Fe-19Ni-9Co

[200-300 grades]

U.S. IV IV (237)

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

austenitic steels VI V (171)

VI (11,127,157)

VII (103) ALUMINUM ALLOYS

Al-2.7Cu-2.2Li0.12ZrU.S. X X (163)

Al-6Cu-0.2Mn-0.1Fe

[5083-0] U.S. II II II II II (19)

Al-0.6Fe-0.2Cu0.1S1

[1100]

U.S. IX IX (107)

Al-4Mg U.S. I

V

II II (35)

[5083-0] IV (185)

COPPER ALLOYS

99 . 99Cu-Ag

[C10100, C10200]

U.S. IX IX (95,107)

99.5Cu-Ag

[C10400

]

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

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

[C15500] 151)

97.6 Cu-Ag-0.38Be

[Cl 7510

]

U.S. X X (83)

SUPERCONDUCTORS

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

*Materials Studies for Magnetic Fusion Energy Applications at Lav Temperatures:

NBSIR 78-884 (I);

NBSIR 82-1667 (V);

NBSIR 86-3050 (IX);

NBSIR 79-1609 (II);

NBSIR 83-1690 (VI);

NBSIR 87-3067 (X).

NBSIR 80-1627 (III);

NBSIR 84-3000 (VII);

NBSIR 81-1645 (IV);

NBSIR 85-3025 (VIII)

;

7

AUSTENITIC STAINLESS STEELS

WITH EMPHASIS ON STRENGTH AT LOW TEMPERATURES*

R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

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 close-packed martensite, stacking- fault energy, carbides,

and sigma phase are described. Recent international development of higher strength, tougher austenitic stainless steels

is summarized. This chapter also includes 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.

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% t ) and low C (<0.3%). Research beginning

in 1904 by Guillet (1904—1914, 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 application of these steels has

To be published in Alloying

,

ASM International, Metals Park, Ohio.

^Unless otherwise specified, wt.% is used.

9
broadened considerably, along with their composition. Only two criteria

determine whether steels are classified as stainless (or rustproof or pitless) : They must be Fe-base elements and they must contain about

HCr or more. This broad classification is divided into six principal

subclasses

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

solid- solution 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. 8 * 9 In 1956,

Keating wrote a user-oriented 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 in

1977 by Peckner and Bernstein12 represents the most comprehensive review

of these alloys; it documents stainless steel melting, fabrication,

metallurgy, corrosion resistance, properties, and applications. Lula

10
recently (1986) revised a general introductory book by Parr and Hanson, 13

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-manganese stainless

steels

.

1

4

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

(a) Fe-Cr-Ni ternary alloys; (b) Fe-Cr-Mn ternary alloys.

11
From the 1950s into the 1980s, several major international conferences have been held on stainless steels in general and specific aspects

of these alloys. Those of which we are aware are:

1. Symposium on cr-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, the some of the first reports of dual -phase

stainless steels were presented.

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

applications; 1977; sponsored by Climax Molybdenum Company and held

in London18

5. Symposium on stainless steel castings; 1980; sponsored by ASTM19

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, Japan20

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

steels. 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 because 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 discussed their solidification, precipitation, and transformation

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

cryogenic properties.

12
STRUCTURES

Stainless steels are complex. They solidify as either an austenitic

(fee) or a ferritic (bcc) structure. At intermediate temperatures, car- bide precipitates, an intermetallic compound (a-phase) , 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.

Austenite

Austenite is naturally the dominant phase of austenite stainless

steels at ambient temperatures. Austenite is sometimes called 7-phase

and is always designated 7 in phase diagrams. Solid- solution alloying

elements, such as Cr, Ni, and Mn, assume normal lattice sites in the Febase, 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 shortrange ordering may be present, but so far the evidence is circumstantial.

Austenite remains paramagnetic on cooling until about 50 K; then it usu- ally becomes antiferromagnetic. The magnetic structure of alloys with

high-Ni content may become mictomagnetic or spin- glass at low temperatures. The Neel temperatures or transition to mixed ferromagnetic states

is strongly dependent on alloying.

Ferrite

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 5 -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 diffusioncontrolled 7-+Q transformation. In wrought alloys, the retained 5 -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, 5 -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 Speich 22

and Schurmann and Brauckman. 23 The eutectic trough extends from near th<-

13
Fe corner (76Fe—lONi—14Cr) to a ternary eutectic (about 8 Fe—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 aus- tenitic stainless steels. The ternary isothermal sections at 1350°C and

1100°C are shown in Fig. 2. At lower temperatures, the S + 7 region increases with attendant decreases in both S and 7 single regions.

The tendency to transform from S -ferrite to 7 is illustrated by

the phase diagrams at constant Fe content but varying Ni and Cr contents

(Fig. 3) . Notice that the minimum Ni content for stable 7 at 70Fe is

about 8 % (22Cr) and at 60Fe, about 14.5% (25.5Cr). In practice, these

limits are adjusted by small additions of C, N, Mn, Si, and, sometimes,

Mo.

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 wt.% alloy contents) at temperatures above 800°C (below 800 e C,

a-phase forms in Fe—Cr alloys, and below 500 B 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 2 Mo and 2 Si. Austenite is encouraged by alloying

Fe with Mn; at 1000°C, the 7 - loop extends to 70Mn.

Chromium

Fig. 2. Isothermal sections at 1350°C (dashed lines) and

1100°C (solid lines) for the Fe—Cr—Ni ternary system.

14
TEMPERATURE,

#C

Fig. 3

CHROMIUM, wt.% CHROMIUM, wt.%

SO 20 10 0

(a) (b)

Binary Fe—Crx—Niy phase diagrams depicting phase dependence on

Cr and Ni at constant Fe content: (a) 30 wt.% Fe; (b) 40 wt.% Fe

CHROMIUM, wt% NICKEL, wt%

Fig. 4. Binary Fe—Cr and Fe—Ni phase diagrams.

15

TEMPERATURE.
Fig. 5. Ferrite portion of Schaeffler diagram depicting lines

that estimate the first solidification of ferrite

during cooling to room temperature after welding.

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

from a Schaeffler26 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 28 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. 29 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, 29 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.

16
Table 1. Ferrite Solidification Equivalent Coefficients

Investigator (Year)

Cr- Equivalent Coefficient

(per wt.%)

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 —

Ni-Equivalent Coefficient

(per wt.%)

Ni Mn C N Cu Co Mn2

DeLong (1956) 1 0.5 30 30 - -

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

Hammar

,

Svensson (1979)

1 0.31 22.5 14.2 1

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 con- sideration 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 (a') and hep

(O martensite phases. The alloy composition affects the chemical freeenergy 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 freeenergy relations between fee and bcc and suggested a relationship for the

free-energy differences between fee and hep structures based on the ear- lier work of Kaufman on binary Fe—Cr and Fe—Ni alloys . 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

at ' -martensite

.

17
Expressions relating the stability of the austenite and the temperatures of transformation during cooling (Tms ) or during deformation (Tm(j)

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 aus- tenite 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 concentrations seemed to be influenced

by the Cr and Ni concentrations.

Table 2. Temperature -Equivalent Coefficients for Calculation of Tms .

Investigator Temperature -Equivalent Coefficient (degrees K per wt. %)

(Year) Base Cr Ni Mn Si C N Mo

Eichelman,

Hull (1953)

1578 -41.7 -61.1 -33.3 -27.8 -1670 -1670

Monkman et al.

(1957)

1455 -36.7 -56.7 -1460 -1460

Andrews (1965) 273 -12.1 -17.7 -30.4 -423 -7.5

Hull (1973) 1755 -47 -59 -54 -37 -2390 -3720 -56*

Self (1986) 794 -14.3 -17.5 -28.9 -37.6 -350 -29. 6t

* Other: -180(Ti)

t Other: -1.19(CrNi) + 23 . 1 (Cr-t-Mo)C

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

,

Th, V, W 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 effects 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 35 applies to lower Cr and Ni concentrations, and thus higher Tms , and his results are included in Table 2 to indicate the disparity between

the high- and low- temperature empirical results. In all studies,

18
specimens of various compositions were prepared and cooled, and their

Tms values were measured. Regression analyses were used to obtain the

empirical dependence of Tms on alloy concentration. Self36 recently

studied a wide range of compositions of weld deposits. For analyses,

he assumed nonideal behavior, and with regression analysis, he derived

a predictive equation containing interactive Cr—Ni and (Cr + Mo)—C terms

(Table 2).

From Tables 1 and 2, it is apparent that the roles of Cr, Mo, and Si are reversed in the stabilization of the bcc structure (ferrite and a'

martensite) between high and low temperatures. Binary Fe-Cr free-energy

derivations 31 * 32 do not suggest this trend. Perhaps the explanation for

the stabilizing effect at lower temperatures is the contribution of these

elements to the increase in austenite flow strength. The affect of Mo on flow 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.

Experiments have been conducted to assess the influence of cold

work (either tensile or compressive) on Tmci- Results are contained

in Table 3. The experimental definition of Tmci 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 higher by a factor of 50; after deformation, Cr and Mn are decidedly less than Ni, and C and N are higher by

a factor of about 10. Careful research to delineate the dependence of

composition on free energy, stacking- fault energy, and deformation

parameters is needed to understand these distinctions.

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—UNi to Fe—lOCr—16Ni ; 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>.y, with (225)~, (112)^, or both habit planes. 41 The laths are restricted within (lll)^ bands, and usually three sets of habit planes

form within the band.

19
NICKEL EQUIVALENT

Fig. 6. Martensite portion of the Schaeffler diagram depicting

lines that estimate the first formation of martensite

during cooling to room temperature after welding.

Constant amounts of lMn, 0.03C, and 0.5Si.

Table 3. Temperature -Equivalent Coefficients for Calculation of Tmci Investigator Temperature -Equivalent Coefficients (degrees K per wt. %)

(Year) Base Cr Ni Mn Si C N Mo Other

Angel

(1954)

686 -14 -9.5 -8.1 -9.2 -46.2 -46..2 -18.5

Hull

(1973)

1655 -23 -59 -41 -20 -777 -315 -24 12 (Co)

Williams et al.

(1976)

686 - 6 -25 -16 +21 -222 -222 -11

20
30.

Iron

Lath

Martensfte

r°

o

o

\ o o

o Surface

Martensite

Plate

Martensite

i I 10 15 20 25

NICKEL, wt.%

30

xi 1 35

Fig. 7. Compositional dependence of a' -martensite

morphology in Fe-Cr-Ni ternary alloys.

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) € || (101V

[H0] 7 || [T210] € || [lll] a

.

These relationships are apparently retained regardless of the manner of e or a' formation.

Transformation during plastic deformation initially results in a'

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, mor.' than one lath forms at each transformation site, and in contrast to the

a' that forms on cooling, all laths have the same habit-plane variant.

The internal defect structure of the a' -phase in Fe-Cr-Ni steels

consists predominantly of dislocations. As Cr is replaced by Ni,

21
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 twin reflections become more diffuse.

Many metallurgists have been concerned with the role of the e transformation: does e act as a precursor transformation ) , or is the

e an accommodation effect (y-»a' , y-*e ) ? Both effects have been observed.

The stainless steels with low stacking- fault energies tend to form e, and

the a' tends to form from e. Higher stacking- fault-energy alloys require

transformation stresses for observable e formation.

Several characteristics of the transformation from austenite to bcc

martensite in stainless steels cause problems in many applications. The

martensitic bcc has a specific volume 1.7% larger than the parent aus- tenite; therefore, these steels expand during transformation. The mar- tensite forms as individual crystals, and the shear stress and volume

expansion associated with transformation disrupt localized regions. For

example, aformation 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 1% of

bcc martensite results in a permeability increase of 0.01. Therefore,

if time -dependent field changes corresponding to material changes on 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

stacking- fault energy leads to wider partial -dislocation separation

(larger fault ribbons) and, hence, more planar dislocation slip structures and reduced cross slip.

Schramm and Reed42 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 y-*e transformation on com- position, because a stacking fault represents a local planar area of e. Rhodes and Thompson43 suggested that the stacking -fault energy values

of the Schramm and Reed analysis were too large, considering additional

weak-beam electron-microscopy data. Weak-beam electron-microscopy mea- surements by Bampton, Jones, and Loretto 44 confirmed that the leastsquare analysis of Schramm and Reed produced stacking -fault energies that

were large. From measurements of individual nodes, they estimated a data

spread of ±25%. Brofman and Ansell 45 added C content to the factors

affecting stacking- fault energy. Dependencies of stacking -fault energy

on composition are contained in Table 4.

22
Table 4. Stacking- Fault Energy at Room Temperature

Investigator Energy Equivalent (mJ/m2 • wt

.

%)

(Year) Base Cr Ni Mn Si C Mo Other

Dulieu, Nutting 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 < 0.325 wt.%, N < 0.21 wt.%) on the stackingfault energy of an Fe—8Cr—lONi 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-beam 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. 48 measured stacking- fault

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

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

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

consistent: the stacking- fault energy increased slightly with N contents

up to about 0.20 wt.%; it decreased substantially for higher N contents.

Carbides

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

in austenitic alloys is very temperature dependent (Fig. 8). 49,50 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—19Cr—14Ni—2Mo) has been represented by the expression51

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

where T is in K. 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. Thi>

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 M 23 C 9 structure

23
LLi QC

D

H

<

CCHI

Q.

HI

CARBON, wt.%

NITROGEN, wt.%

Fig. 8. Solubility limits for carbon49 and nitrogen50 as a function of temperature for various alloys.

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, 52 Fig. 9

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

data of Sakamoto et al., 53 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 prefer to precipitate at ferrite—austenite interfaces, followed by grain boundaries, noncoherent and coherent twin boundaries, and,

finally, within grains. This is depicted schematically in Fig. 10. 54,55

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. 1

2

24
Fig. 9. Solubility limits of nitrogen in Fe—Cr, Fe—Ni, and Fe—Mn

binary alloys at 1473 K

52 and in Fe—15Ni—Cr—Mn alloy liquids. 53

The M23 C 6 carbide has a complex fee structure with 92 metallic atoms

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

10.64 A . 12 But there is only minor mismatch between the (111) M23 C 6 and

{111} austenite, estimated at 1.3% by Lewis and Hatterslay

.

5

6

Thus, it

is not surprising that the common plane of the interface between the two

structures has been identified as (111). 65,56 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

.

Carbide precipitates normally have a sheet- like morphology at lower

temperatures (480—730°C) , dendritic shapes at intermediate temperatures

(600—875°C), and are particles at the higher temperatures (>850°C)

.

5

5

Higher temperatures, as well as longer aging times, promote coarsened

structures

.

25
*

llT QC

o

<

CO

LUQ_

2

LU

tFig. 10. Time—temperature carbide precipitation kinetics,

with specific sites identified. 54 > 55

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 6 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 Novak25 and Lai. 57

26
Sigma Phase

Notice the Fe-Cr part of the phase diagram in Fig. 4: at Cr con- tents greater than about 15%, the a-phase may form when the alloy is

exposed to temperatures between 400 to 800°C. The a-phase extends into

the Fe—Cr—Ni ternary diagram, as indicated by Fig. 3. Alloys rich in Cr

and lean in Ni tend toward a-phase formation at lower temperatures.

Maehara et al. 58 studied the effects of alloy additions on the amounts

and aging times for a-phase formation. They found that Ni decreases the

amount of a-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 a-phase in Fe—6 . 5Ni—2 . 8Mo and Fe—25Cr—6 . 5Ni-base

alloys, respectively.

In Fe—Cr, the highest temperature of a-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

Speich22 and the 802 °C and 648 °C isotherms of Talbot and Furman. 59 Like

carbides, the a-phase forms preferentially at high-energy surfaces.

Also, when ferrite is present, there is a strong tendency for a-phase

to form initially within the ferrite or at the ferrite side of the j—6

boundary.

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

aQ lattice parameter range of -8.29 to 9.21 A, and c 0 lattice parameter

range of —4.60 to 4.78 A. 12 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 Novak25 and Lai. 57

27
PROPERTIES

Stress—Strain Characteristics

The AISI 300 series stainless steels used in cryogenic applications

range from metastable to stable austenites. Alloys such as 304, con- taining 18Cr and 8Ni, are metastable, whereas 310, containing 26Cr and

20Ni is stable with respect to martensitic transformation. The stressstrain 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. 11b) 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.

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 single -crystal shear- stress—strain curves. Stage I represents the microstrain and early microstrain behavior. The formation

of a' -martensite is not thought to occur in this range; stacking- fault

clusters or e -martensite , or both, are most likely to complement dislocation interactions to reduce the rate of work hardening.

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

e-martensite formation and the formation of a' -martensite laths at cross- slip intersections. Suzuki et al. 60 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 is 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 to being associated with a linear, high rate of work hardening

in stage III—is not clear. Perhaps in stage III essentially all active

28
STRESS,

MPa

STRAIN

(a)

250

200

150

100

50

t/> JC

CO* CO

cu

cn

i—CO

(b)

Fig. 11. Low- temperature stress-strain curves of

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

(b) a metastable austenitic alloy 304 (Fe-19Cr-9Ni )

.

29
Elastic

Plastic

Stages

Fig. 12. Schematic depiction and characterization of three stages

associated with the stress-strain curves for polycrystalline

metastable austenitic stainless steels at low temperatures.

cross-slip sites have been transformed to a-martensite , and subsequent

a' -martensite formation occurs within active glide bands. Such a '

martensite formation would not be expected to promote stacking- fault

glide. There is a hint from the x-ray data of Reed and Gunter 81 that

the following sequence may occur: A maximum value of the volume concen- tration of e -martensite occurs at about the transition from stage II to

stage III. There is apparently no e -martensite formation during stage

III; the amount of e -martensite decreases with strain, presumably transforming to a' -martensite

.

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

but 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 e -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. 82 The yield strength, <7y (flow strength at 0.002

plastic strain) is within stage I. As shown in Fig. 13, Oy is linearly

dependent on Mn content, increasing with increasing Mn. The linear dependence matches that of ajj, 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

30
1.6

1.4 -

1.2 -

o

Q_O i.oh

^ 0.8

0.6

0.4 1

1 2 3 4 5 6 7

MANGANESE, wt. %

Fig. 13. Dependence of stages I and II strength

parameters on Mn and N content at 4 K.

8

parameter nor the shear modulus significantly (see later discussion)

,

therefore traditional solid- solution strengthening should not result from

Mn addition. But Mn increases the stacking- fault energy and suppresses

a'- and e-martensite transformation; these phase transformations influ- ence the stress and strain parameters. Therefore, the increases in

strength that occur when Mn is added are due to the contribution of Mn to

austenite stability. Figure 13 shows that an addition of 0.1N increases

ay about 290 MPa. • Nitrogen has the same strengthening effect in stable

alloys, which implies that it has little effect on hep or bcc stability

within the composition range 0.1 to 0.2 wt.%.

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

rates, ajjj, 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.

31
Fig. 14.

O

CL

O

x

o

b

X5

VL> "O

Dependence of stage

III work-hardening

rate on Mn and N

contents at 4 K. MANGANESE, wt. %

Fig. 15. Dependence of stage-XII work-hardening rate on the normalized

amount of a-martensite formed during uniform elongation at 4 K.

32
2.4

Fig. 16. Correlation between normalized a ' -martensite formation

in stage III with alloying, as characterized by the

Tmd calculations of Williams et al. 37

We have performed numerous correlations between the normalized

amount of strain- induced a' -martensite at final elongation at 4 K and

the alloy composition. Our best correlation for this normalized straininduced martensite was obtained by using the Williams et al. 38 regression

analysis for Tmcj. In Fig. 16 their expression is plotted versus a ' /%

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

stress—strain curves of metastable austenites during and strain- induced

a'- and e-martensite transformations is affected by alloying.

33
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. 63-66 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 a yield strength of 1200 MPa and a toughness of 200 MPa-m^

has now been completed. 67—69 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 " 83 ' In Fig. 17, the effect of C + N on yield

strength at 4 K for an Fe—18Cr—lONi steel is presented. 72 The break in

b*

CO

CO

111 CL

I—CO

o

UJ>

Fig. 17. Dependence of yield strength on C + N contents in Fe—18Cr—lONi

austenitic stainless steels at low temperatures.

34
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 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 yield

strength increased 237 MPa per 0.1N; 77 an earlier analysis of 30 data

points for 304LN-type alloys at 4 K found that yield strength increased

319 MPa per O.IN. 70 Strengthening from C is about half that from N at

4 K. 70 Thus, N addition strengthens considerably more than equivalent

C addition 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. 64,65,70-75,77-83 Reed and Simon70 ’ 77 found a smaller

standard deviation from regression analyses by using the power of one

to express the N- content dependence rather than a 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 yield

strength on C + N shown in Fig. 17 also supports this assessment. This

dependence of yield strength 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 yield strength 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. 74,87-83 There

is evidence that at higher Ni contents the strengthening effects of Mo

are reduced. 89 Manganese, when not affecting austenite stability, increases yield strength by about 10 MPa per wt.% at 4 k. 53 ’ 73 ’ 75 ’ 89 ’ 90

Nickel does not contribute directly to solid- solution strengthening,

but it does affect strength through its strong contribution to austenite

stability. The temperature dependence of the solid-solution strengthening additions has not been assessed.

In Table 5, the strengthening contributions from substitutional and

interstitial alloying additions in fee elements 92-97 and in austenitic

stainless steels 62-91 are listed in terms of a fraction of the shear

modulus. This refers to the increase of yield strength 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.%.

35
Table 5. Strengthening Contributions in fee Alloys

Solvent Solute Temperature Amount

(Act/100%)

Ag Ai 4 G/30

Sb 294 G/100

76 G/30

Ni Mo 295 G/45

C 4 G/7

295 G/30

Cu Ni 76 G/120

Au 76 G/100

Pb A * 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 °C

Adding C to Ni95 o’r 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 Goldman92

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, 84-86 the shear

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

with mobile dislocations is

r = A[CF

2/3] [f0V3

], (2)

where Cp is the concentration of a solute species, and A is considered

to be a constant for this discussion. The maximum interaction force is

given by

36
r, 2 + a

2 62 f1

2

, (3) o where the modulus misfit is

. 1 dG

5 G dc*

(4)

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

misfit

,

S ! 4s

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 Austin98 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, u 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

1 da 1 dE 1 dG 1 dB 1 du V(X)

a dc E dc G dc B dc V dc V ( Fe

)

295 K 295 K 4 K 295 K 4 K 295 K 4 K 295 K 4 K

cN

0.218

0.240

-0.78 0.61 -0.76 0.81 -0.93 -0.52 -0.11 -0.90 0.34

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 oo lH H -0.38 0.79 1.30 1.22 1.13 1.32

Cr 0.075 — — — — — — — — 1.02

37
At room temperature both C and N show a large size 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 for C, although the atom size of N is

less than that of C. The divergence of E and G trends from B and u 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 changes in sign of

the Young's and shear moduli parameters of N and C on cooling from 295 to

4 K imply 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 with 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 than 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 sizemisfit 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, leads to the con- sideration of interstitial clustering.

Galligan and Goldman92 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 en- vision 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 very temperature

dependent

.

JL 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

*fg - r) 2 + a2 S

2

, where a is assumed to equal 20

38
SUMMARY

Austenitic stainless steel structures and strengths have been

reviewed, with emphasis 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 hcp-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 to more

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 best alloying element for

increasing strength at low temperatures.

ACKNOWLEDGMENT

This work was partially funded by the Office of Fusion Energy,

U.S. Department of Energy.

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Cryogenic in Engineering — Materials , vol. 28, Eds. R. P. Reed and

A. F. Clark, Plenum Press, New York (1982), pp. 83—92.

73. T. Horiuchi, R. Ogawa, M. Shimada, S. Tone, M. Yamaga, and

Y. Kasamatsu, "Mechanical Properties of High Manganese Steels at

Cryogenic Temperatures," in: Advances in Crvogenit Engineering —

Materials . vol. 28, Eds. R. P. Reed and A. F. Clark, Plenum Press,

New York (1982), pp. 93-103.

74. S. Yamamoto, N. Yamagami, and C. Ouchi, "Effect of Metallurgical

Variables on Strength and Toughness of Mn—Cr and Ni-Cr Stainless

Steels at 4.2K," in: Advances in Cryogenic Engineering — Materials , vol. 32, Eds. R. P. Reed and A. F. Clark, Plenum Press, New York

(1986), pp. 57-64.

75. T. Horiuchi, R. Ogawa, and M. Shimada, "Cryogenic Fe—Mn Austenitic

Steels," in: Advances in Cryogenic Engineering — Materials . vol. 32, Eds. R. P. Reed and A. F. Clark, Plenum Press, New York

(1986), pp. 33-42.

76. R. L. Tobler, D. H. Beckman, and R. P. Reed, "Factors Influencing

the Low Temperature Dependence of Yielding in AISI 316 Stainless

Steels," in: Austenitic Steels at Low Temperatures . Eds. R. P. Reed and T. Horiuchi, Plenum Press, New York (1983), pp . 135—157.

43
77. N. J. Simon and R. P. Reed, "Design of 316LN-Type Alloys," in:

Advances in Cryogenic Engineering — Materials , vol. 34, Eds. A. F. Clark and R. P. Reed, Plenum Press, New York (1988), pp. 165—172.

78. R. P. Reed, P. T. Purtscher, and K. A. Yushchenko, "Nickel and

Nitrogen Alloying Effects on the Strength and Toughness of Austenitic Stainless at 4 K, in: Advances in Cryogenic Engineering -

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

New York (1986), pp. 33—42.

79. T. Sakamoto, Y. Nakagawa, and I. Yamauchi." Effect of Mn on the

Cryogenic Properties of High Nitrogen Austenitic Stainless Steels,"

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. Masumoto, K. Suemune , H. Nakajima, and S. Shimimoto, "Development

of High-Strength, High -Manganese Steels for Cryogenic Use,” in:

Advances in Cryogenic Engineering — Materials , 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," Chapter 15 in:

Dislocations in Solids . 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.

35, 613-622 (1977).

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

Metall . 20, 917-927 (1972).

87. K. Suemune, K. Sugino, H. Masumoto, 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.

44
89 . P. T. Purtscher, R. P. Walsh, and R. P. Reed, "EffecC 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 (1988),

pp. 191-198.

90. R. Muira, H. Nakajima, Y. Takahashi, and K. Yoshida, "32Mn—7Cr

Austenitic 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. Ishikawa, 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 SolidSolution Strengthening," in: Strength of Metals and Alloys,

Proceedings of the Fifth International Conference . Pt. II, Peramon,

New York (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. 33, 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. den Otter and A. Van den Beukel, "Flow Stress and Activation

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

.

Status Solidi(a) 55, (1979), pp. 785-792.

97. J. H. Tregilgas and J. M. Galligan, "Hardening from Metal Interstitials in a Face-Centered Cubic Lattice—Pb—Ag," Scr. Metall. 9,

1225-1227 (1975).

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

Bureau of Standards, Boulder, Colorado (1987).

45

LOW- TEMPERATURE PROPERTIES OF

HIGH -MANGANESE AUSTENITIC STEELS^

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

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

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.

INTRODUCTION

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 lowtemperature applications, such as those in magnetic -fusion- energy

reactors and high-energy-physics 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 1 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) aus- tenitic 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

•||i To be published in High -Manganese Austenitic Steels, ASM International,

Metals Park, Ohio.

47
HIGH -MANGANESE AUSTENITIC STEELS

Fig. 1. Portion of the Fe—Mn—Cr ternary diagram with compositions

of recently developed high-Mn austenitic steels.

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

high-Mn austenitic steels at cryogenic temperatures. The most significant design parameters—strength, toughness, and physical properties

are emphasized.

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 e (hep) transformations nor the strain- induced a' and e' transformations

are significantly affected. Additions of Cr greater than approximately

48
15% result in high- temperature <5 ferrite or a phase (at high Mn con- tents). Alloying with relatively small amounts of Ni, C, or N is needed

to suppress these high- temperature phases that are deleterious to fracture. Figure 2 also shows 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

c -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 3 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 e nor e' was detected at Mn

contents higher than 25%.

Additions of Ni and C to Fe—Mn alloys enhance austenite ( 7 ) stability. The Fe-rich portion of the Fe—Mn—Ni ternary diagram5 containing

both thermally and strain- induced phases at 76 K is shown in Fig. 3.

Fig. 2. Portion of the Fe—Mn—Cr ternary diagram illustrating allov dependence of phases formed during deformation or cooling to 4 K

.

:

*4.11 compositions in this paper are in wt.% unless otherwise specified.

49
50,

Fig. 3. Portion of the Fe—Mn—Ni ternary diagram illustrating alloy

dependence of phases formed during deformation or cooling

to 76 K. 5

MANGANESE, wt.%

Fig. 4. Portion of the Fe—Mn—C ternary diagram illustrating alloy

dependence of phases formed during deformation or cooling

to 300 K. 6

50
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 e' 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

e martensite, and transformation to a' martensite.

Slip . 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 , 7 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' -martensite 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 longrange 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 a' 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 t

' 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 . 8 of Fe—31Mn and Fe—36Mn-alloy flow strengths i.:; 1 v that twinning does not substantially alter the tensile behavior, as slip

does

.

51
Fig. 5. Deformation mechanisms in Fe—Mn—(Cr or C) alloys

as a function of temperature and Mn content. 7

Transformation to e' Martensite . Transformation to e' martensite,

rather than slip or twinning, occurs at lower temperatures and Mn con- tents (Fig. 5). Formation of e' martensite during deformation increases

the rate of work hardening. 8 The platelets of e' on {111}^ are thought

to segment the austenite grains, considerably obstructing the dislocation

motion through effective reduction in grain size. This effect of e' formation on stress—strain behavior is illustrated in Fig. 7. In the ab- sence of a' formation, strain- induced e' 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, 9 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>y axial direction corresponding to the common direction

of { 1 1 1 } -y intersection. These laths effectively reduce the flow stress

52
Fig. 6. Photomicrograph of Fe—22Mn—13Cr—5Ni—0 . 22N austenitic steel

deformed in tension to 10% elongation at 76 K.

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 a' forms from e' . 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

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 a' , 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). 11 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

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

53
STRAIN

Fig. 7. Stress—strain and associated martensite- transformation

characteristics for austenitic steels at low temperatures.

54
atom produces larger changes in the lattice parameter (a) : at 295 K,

(1/a) (da/dc) - 0.218 for C and 0.240 for N !

,

1

1

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 dependences of yield strength (cry, in MPa) on C and N content

([C] and [N] , in wt.%) and grain size (d, in /im) were analyzed using the

expression

oy - a 0 + a[C] + b [N] + kd -4

, (1)

where cr 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

2 + a2 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 dependences of the C and N

strengthening contributions are larger than that of a Q , implying shorter

range stress fields from the interstitials. (2) The ratio of these con- tributions 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 austenites. Furthermore, in the Fe—Cr—Ni alloys,

Table 1. Regression Analysis of C, N, and Grain-Size Contributions

to Tensile Yield Strength in Eq. 1

Tempera- Number Coefficients Standard Deviation

ture of Data (MPa/wt.%) (MPa/wt . %

)

(K) Point o Q a b k Oy ac

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

55
Fig. 8. Regression analyses results for the dependence of yield strength

on C and N content and grain size at low temperatures.

56
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 analyses 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, cr 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. Lula29 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, 9 which leads to the suspicion

that N is positioned differently in the lattice than C, possibly in

pairs 9 or ordered within short range. 30

Grain Size Dependence

The published data 1

5

"

1

7

* 20 * 26 » 3

1 _33 on the effects of grain size on tensile yield strength, using the Hall—Petch relationship, are included

in Fig. 9. Here,

ay ” ao

+ kd'^,

where a 0

' includes all grain- size -independent contributions.

(3)

Table 2. Nitrogen and Carbon Contributions to Tensile Yield Strength

Source Alloy 4 K

Nitrogen

76 K

(MPa/wt.%)

295 K 4 K

Carbon

76 K

(MPa/vt.%)

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( 28

) Fe-Ni-Cr-Mo 2780 1690

Did not include a grain-size term in the analysis

57
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—Ni alloys 28 is 1090 MPa/pm"^; for Fe—Cr—Ni—Mo

alloys, 810 MPa/^m'^. 3

4

The lower average values for these alloys are outside the data spread of individual Fe—Mn—Cr-alloy studies. The

increased grain-boundary contribution of the Fe—Mn—Cr alloys implies

increased alloy segregation at their grain boundaries, presumably by

C and N.

cr

y = cr0 + kef

1/2

Fig. 9. Summary of literature data on the Hall—Petch constants of

high-Mn austenitic steels at low temperatures. Also included

are the k results of the regression analyses of this study.

58
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. 3S 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.

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 0 "exp(-AT) , (4)

where ct 0

" is the yield strength at absolute zero and A is a constant, the

slope of log cry 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

Fig. 10. Temperature dependence of austenitic alloy yield strengths

at low temperatures

.

59
signifies a change of the deformation mechanism. Many of these alloys

are completely stable: there is no strain- induced martensite. The ma- croscopic 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 ’ 15 Recent studies 4 ’ 15 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.

At low temperatures, Fe—Mn—Cr and Fe—Mn—Cr—Ni alloys with high N con- tents or low austenite stabilities exhibit faceting of the stackingfault clusters on {lll}.y in impact or fracture toughness tests. 4 * 36 " 38

This brittle- looking faceting is similar to that reported by Tobler and

Meyn39 as slip-band cracking on compact- tension specimen fracture sur- faces in high-N alloys. These facets usually have other {111}^ traces on their surfaces (Fig. 11a) . The fracture surfaces following fracturetoughness tests at low temperatures are seldom dominated by these {111}^

facets; elongated, shallow dimples are the prevalent morphological structure (Fig. 11a). At higher temperatures or in more stable structures,

the faceting is reduced and the elongated dimples are dominant on the

fracture surfaces (Fig. lib). 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—2 2Mn— (13—15)Cr—(1—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.

60
Fig. 11. (a) Fracture surface of a compact- tension specimen of Fe—8Mn—

20Cr—7Ni—0 . 37N alloy, tested at 4 K, displaying {111) 7 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.

61
YIELD STRENGTH, cr

y

, MPa

Fig. 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

.

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

Kic(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 (Jj c measurements are about 10 to 20% lower

in this orientation than in the LT orientation) . The tensile yieldstrength data in Fig. 12 were measured in the transverse direction. The

individual data points portray our measurements of high-Mn austenites.

62
The several data points of the solution- treated alloys 18Mn—16Cr—5Ni—0 . 2N

and 22Mn—13Cr—5Ni—0 . 2N reflect test and material variability. The location of these high-Mn-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; low-yield- strength alloys have larger dimples.

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

.

Fig. 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-22Cr13Ni—5Mn, (b) Fe-21Cr-6Ni-9Mn, and (c) Fe-18Cr-3Ni-l 3Mn

.

63
The transition to an antiferromagnetic state produces a softening

of the elastic properties. This softening is relatively large, as illustrated by Ledbetter's data43 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%.

The magnetic transition also strongly affects thermal expansion.

Below the Neel temperature (Tjj) , there is positive magnetostriction in

the antiferromagnetic state, which reduces the thermal -expans ion coefficient. Even at temperatures up to 100 K above the transition temperature, there is limited magnetic ordering that lowers the coefficient.

Richter and Pepperhoff44 have measured the thermal -expans ion characteristics of a number of Fe—Mn base alloys at low temperatures. They observed that C and Cr additions tend to lower and reduce the volume

magnetostriction (Fig. 15).

Fig. 14. Elastic properties for an Fe—26Mn—5Cr alloy at low temperatures

[Ledbetter39

]

.

64
100

Io

0

crCO

t

1

s

80 -

60

40 —

20

Fe-31 Mn-xC

Fe-25 Mn-.OlC-xCr

Fe—xMn

x Fe—Mn

o Fe—31Mn—

C

Fe-25 Mn-.OIC-Cr

i I

ALLOYING ELEMENT, wt % (0.01 % for C)

Fig. 15. Effects of alloying on Neel temperature and volume

magnetostriction of Fe—Mn-base alloys. 40

In the Fe—Mn binary system, Mn additions of 32 to 50% increase T^,

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 .

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 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.

65
0

o

8 o.i o

o<ccoo

0.2

0.3

3

04

at Aluminum

100 200

TEMPERATURE, K

300

Fig. 16 Percent contraction of various face-centered

cubic structural alloys on cooling.

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 { 1 1 1 } -y is observed in

lower stacking- fault-energy alloys. Binary Fe-Mn alloys are more sus- ceptible 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

.

ACKNOWLEDGMENT

This work was supported, in part, by the Office of Fusion Energy,

U.S. Department of Energy.

66
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21. Tone, S., M. Hiromatsu, J. Numata, T. Horiuchi, H. Nakajima, and

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Centre, Paris, France (1986).

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587-95 (1981).

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Engineering — Materials," vol. 28, pp. 115—25, Plenum Press, New

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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).

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Engineering — Materials," vol. 34, pp. 299—308, Plenum Press, New

York (1988)

.

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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).

68
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. , Physica 1 19B . 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).

69

DESIGN OF 316LN-TYPE ALLOYS*

N. J. Simon and R. P. Reed

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

For structural use in 4-K superconducting-magnet applications, alloy 316LN offers higher strength than 304LN with

equivalent toughness. Nitrogen additions permit the attainment

of high yield- strength levels with little added cost. To meet

the goal of combined high yield strength and high fracture

toughness at 4 K, careful consideration of alloy design parameters is necessary. 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, 316LN, and 316LHN

alloys. This data matrix includes more extensive numerical

data on alloying, refining, and processing parameters than have

previously been available. The alloys were obtained from suppliers in Japan, Europe, and the United States. Predictive

equations are presented for the yield strength at 4 K as a function of N and Mo content and grain size. Predictive equations for the 4-K fracture toughness, Kjc (J), of these alloys

are derived from a matrix of 45 measurements. The inverse re- lationship of Kj c

(J) to yield strength established from earlier

NBS data at 4 K had allowed an approximate prediction of Kj c (J)

if the yield strength was known. By including the effects of

Ni content and inclusion spacing in the regression equations,

Kic(J) at 4 K can be predicted with higher certainty.

A condensed version of this paper was published in Advances in Cryogenic

Engineering — Materials, vol. 34, Plenum, New York (1988), pp . 165-172.

71
INTRODUCTION

Fusion energy reactors with 4-K superconducting magnets will re- quire cryogenic structural alloys with improved combinations of yield

strength and fracture toughness at 4 K. These structural alloys are required to withstand large magnetic forces in a superconducting magnet

system of restricted dimensions, and thus, high strength is needed. In

addition, because most structures are complex and have regions of high

stress concentration, good fracture toughness is required. However,

fracture toughness correlates inversely with yield strength at 4 K, 1 * 2 and the quantitative dependence of 4-K strength and toughness upon other

parameters has been undetermined. Recently, we deduced predictive equations for strength and toughness at 4 K in which parameters of alloying,

refining, and processing were expressed quantitatively. 3 * 4 Since then,

testing of a series of 316LN and 316LHN alloys at NBS has greatly ex- panded the data base of 4-K measurements of strength and toughness. In

this paper, using the enlarged data base, we re-examined and expanded the

predictive equations to include additional parameters.

The strength of austenitic steels at 4 K depends quite strongly on

N content: each 0.1 wt.% N raises the tensile yield strength ( ay, 0.2%

offset) by approximately 300 MPa. 5 Other elements, such as C, Ni, Mn,

and Cr, contribute much less to the low- temperature strength. The effect

of Mo, however, is sizeable; Oy measurements on 316 -type alloys with Mo

content above the AISI specifications (2 to 3 wt.%) are included in the

present regression analyses to permit a quantitative determination. The

yield strength is also strongly dependent upon grain size; an increase

of grain size from 20 to 100 pm results in a decrease of Oy of about

90 MPa. 8 Thus, yield strength can be significantly increased through

alloying and thermomechanical processing.

Toughness is adversely affected by increased strength: the elastic

stress -intensity factor, Kj c (J),* has a negative linear dependence upon

ay at 4 K. 1 * 2 Toughness also is adversely affected by the presence of

inclusions, 6 ' 7 which act as initiation sites for microvoid formation.

The addition of some alloying elements, such as Ni, has previously been

shown to increase the low- temperature toughness of austenitic steels when

strength and inclusion content are held constant. 8

NBS measurements of Kj c

(J) at 4 K on 32 specimens that meet the

316LN or 316LHN specifications range from 147 to 390 MPa-m*4

. The arithmetic mean is 227 MPa-m^, the standard deviation (SD) is 74 MPa-m^, and

the coefficient of variation, defined as the ratio of the SD to the mean,

is 32%. Such large scatter in mechanical -property data is typical at

4 K: a much smaller variation is normal at room temperature. 9 To

'jfc* - Kj c is obtained from measurement of the elastic—plastic critical

J-integral (Jjc ) and calculated from Kjc

(J) - (Jjc °E)^ where E is

Young ' s modulus

.

72
analyze this variability, an expanded set of 4-K strength and fracture toughness data on 316-type alloys was assembled, combining recent

316LN and 316LHN test data with earlier NBS 316- type-alloy measurements.

This expanded data set, which includes measurements of grain size and

inclusion spacing for each specimen, permitted the derivation of predictive equations that expressed the effects of grain size and composition

on strength and the effects of strength, alloying, and inclusions on

toughness. The predictive equation obtained for Kjc

(J) has a standard

deviation equivalent to a coefficient of variation of 14% , which is a considerable improvement over the coefficient of 32% quoted above.

(The predictive equation was obtained from the expanded data set of 45

measurements, which has a coefficient of variation of 36%. ) Alloy design

(in terms of strength and toughness) is now feasible, within statistical

limits of uncertainty. Of course, at this time, only research data are generally available. Before large-scale use of one of these alloys in a low- temperature structural application, additional heat-by-heat qualification testing is imperative.

PROCEDURE

A series of specimens from both laboratory and commercial heats

of AISI 316-type alloys were tested. The alloys were prepared in the

United States, Japan, and Europe. 10 " 14 Measurements were made of fracture toughness [Kj c (J)] (predominantly TL orientation), tensile yield

strength, and other tensile properties at 4 K. Grain size, hardness,

and chemical composition were also determined, and the inclusion counts

were obtained, as described by Reed et al. 6 Only measurements made at

NBS were used in the data matrices , because previous experience has indicated that lab -to -lab variations in test results are large enough to

obscure the dependence of cryogenic properties upon some parameters.

There are 45 complete rows in the Oy x Kjc

(J) matrix for AISI 316-

type alloys. When several <jy measurements on a particular material were

available, they were averaged and the mean Oy was paired with the Kj c (J)

values. The additional tensile data without corresponding Kj c

(J) mea- surements were used in the regression analyses of Oy\ a total of 99

values were available. Table I shows the data used in the regression

analyses and the reference symbols. (A listing of Gy in table I without

grain size, d, denotes a value that was using in the Kj c analysis only.)

Table II presents the chemical analyses of the alloys studied. Nitrogen content ranges from 0.03 to 0.20 wt.%; nickel content, from 10.5 to

14.2 wt.%; and Mo content, from 2.0 to 4.0 wt.%. Variation of Mo content

outside the AISI 316 specifications permitted a more accurate determination of its effect on ay. Grain size varies from 11 to 242 /im in the

matrix. Inclusion spacing, A, ranges from 0.044 to 0.11 mm (A - n'*1

, where n - the number of inclusions per mm2 ). The Kj c (J) values obtained

from the relationship Kj c

(J) — [Jj c *E/(l — I/ 2 )]

1*, where v is Poisson's

ratio, were adjusted in accord with the formulation Kj c (J) - (Jj C 'E)\

All data points of the matrices were weighted equally in the linear

least-squares regression analyses used to fit the series of equations

described below.

73
Table I. Combined <jy and Kj c Matrices

,

AISI 316 -Type Alloys

Alloy &

N,

wt.%

d,nm KIc,

MPa -in 4 Ni,wt.%

Mo,

wt.%

A,mm

A

10 937 0.16 63 13.7 2.70 —

933 0.16 63 — 13.7 2.70 —

B

10 942 0.16 70 — 11.4 2.50 —

924 0.16 91 318 11.4 2.50 0.109

918 0.16 129 335 11.4 2.50 0.109

1076 0.16 13 214 11.4 2.50 0.129

1096 0.16 13 218 11.4 2.50 0.100

C*io 832 0.10 — 390 13.9 2.20 0.096

834 0.10 31 360 13.9 2.20 0.096

831 0.10 31 380 13.9 2.20 0.096

788 0.10 104 — 13.9 2.20 —

786 0.10 103 — 13.9 2.20 —

>

783 0.10 100 —13.9 2.20 —

783 0.10 100 —

>

13.9 2.20 —

111 0.10 100 13.9 2.20 —

779 0.10 100 — 13.9 2.20 —

790 0.10 100 — 13.9 2.20 —

780 0.10 100 — 13.9 2.20 —

790 0.10 93 — 13.9 2.20 —

772 0.10 100 —13.9 2.20 —

779 0.10 100 — 13.9 2.20 —

779 0.10 100 — 13.9 2.20 —

111 0.10 100 — 13.9 2.20 —

753 0.10 100 — 13.9 2.20 —

D

10 1056 0.13 147 10.9 2.20 0.080

885 0.13 20 187 10.9 2.20 0.083

869 0.13 30 — 10.9 2.20 —

842 0.13 50 — 10.9 2.20 —

935 0.13 11 177 10.9 2.20 0.081

E* 11 1020 0.17 105 200 13.9 2.50 0.057

992 0.17 105 214 13.9 2.50 0.057

p*i 1,12 718 0.03 65 266 13.5 2.30 0.054

704 0.03 65 254 13.5 2.30 0.054

711 0.03 — 243 13.5 2.30 0.054

663 0.03 67 — 13.5 2.30 —

656 0.03 67 — 13.5 2.30 —

* commercial heat

74
Table I. Combined Oy and Kj c Matrices, AISI 316-Type Alloys (continued)

A11°? &

N,

wt.%

d,

fim

KIc,

MPa -in 4 Ni,

wt . %

. Mo,

wt.%

A,

mm

G*i 3 ,i4 55 0 0.03 64 406 12.2 2.16 0.092

690 0.03 22 — 12.2 2.16 —

661 0.03 28 — 12.2 2.16 —

640 0.03 32 — 12.2 2.16 —

553 0.03 49 — 12.2 2.16 —

540 0.03 55 — 12.2 2.16 —

562 0.03 58 — 12.2 2.16 —

535 0.03 61 — 12.2 2.16 —

535 0.03 81 — 12.2 2.16 —

530 0.03 118 — 12.2 2.16 —

505 0.03 138 — 12.2 2.16 —

H* 10 882 0.16 85 203 10.5 2.12 0.044

854 0.14 64 195 10.5 2.12 0.044

862 0.16 67 — 10.5 2.12 —

860 0.14 70 181 10.5 2.12 0.052

848 0.15 70 159 10.5 2.12 0.052

845 0.15 63 — 10.5 2.12 —

856 0.15 66 — 10.5 2.12 —

851 0.15 67 — 10.5 2.12 —

852 0.15 66 — 10.5 2.12 —

847 0.15 73 — 10.5 2.12 —

874 0.15 62 — 10.5 2.12 —

922 0.15 73 — 10.5 2.12 —

857 0.15 64 — 10.5 2.12 —

886 0.15 63 — 10.5 2.12 —

836 0.15 58 — 10.5 2.12 —

856 0.15 60 — 10.5 2.12 —

877 0.15 64 — 10.5 2.12 —

927 0.17 75 — 10.5 2.12 —

859 0.15 - 51 — 10.5 2.12 —

922 0.17 75 — 10.5 2.12 —

919 0.17 75 — 10.5 2.12 —

821 0.15 48 — 10.5 2.12 —

934 0.17 57 — 10.5 2.12 —

927 0.17 87 — 10.5 2.12 —

868 0.15 56 — 10.5 2.12 —

919 0.17 78 — 10.5 2.12 —

835 0.15 48 — 10.5 2.12 —

945 0.17 75 — 10.5 2.12 —

927 0.17 78 — 10.5 2.12 —

859 0.15 51 — 10.5 2.12 —

905 0.17 75 — 10.5 2.12 —

* commercial heat

75
Table I. Combined a-y and Kj c Matrices, AISI 316-Type Alloys (continued)

Alloy <7y , MPa

N,

wt.%

d,

/xm

Klc,

MPa • in* Ni,

wt.%

Mo,

wt.%

A,mm

I 10 1049 0.18 34 163 12.0 2.06 0.055

1014 0.18 28 172 12.0 2.06 0.055

945 0.18 100 — 12.0 2.06 —

1109 0.19 34 157 11.8 3.00 0.053

1119 0.19 35 148 11.8 3.00 0.053

1096 0.19 70 —11.8 3.00 —

1222 0.19 — 98 12.1 4.08 0.057

1181 0.19 —103 12.1 4.08 0.057

1068 0.19 75 12.1 4.08 —

1063 0.17 —

188 13.8 2.05 0.058

1060 0.17 — 202 13.8 2.05 0.058

1129 0.18 — 167 14.0 2.98 0.057

1109 0.18 — 181 14.0 2.98 0.057

994 0.18 100 14.0 2.98 —

1162 0.19 — 98 14.2 4.02 0.051

931 0.17 47 211 13.8 2.05 0.063

934 0.17 57 226 13.8 2.05 0.063

863 0.17 127 241 13.8 2.05 0.063

894 0.17 68 201 13.8 2.05 0.063

903 0.17 94 213 13.8 2.05 0.063

905 0.17 118 210 13.8 2.05 0.063

942 0.19 —149 14.2 4.02 0.040

968 0.19 —

165 14.2 4.02 0.040

915 0.19 —

145 14.2 4.02 0.040

1013 0.19 78 193 14.2 4.02 0.040

1007 0.19 118 170 14.2 4.02 0.040

981 0.19 187 195 14.2 4.02 0.040

j* 10 937 0.15 66 —

_

10.5 2.12 —

916 0.15 58 — 10.5 2.12 —

891 0.16 51 __ 10.5 2.12 —

873 0.15 58 — 10.5 2.12 —

908 0.15 58 10.5 2.12 —

883 0.15 58 —

10.5 2.12 —

Kti° 923 0.18 242 350 13.9 2.70 0.107

923 0.18 — 360 13.9 2.70 0.107

^0 commercial heat

t electroslag remelted

76
Table II. Chemical Composition of AISI 316-Type Alloys

Alloy Cr Ni Mn Mo C N P S Si

A

10 16.7 13.7 1.26 2.7 0.014 0.16 0.016 0.010 0.41

B

10 18.5 11.4 1.40 2.5 0 . 006-

0.010

0.16 0.002 0 . 0002-

0.001

0.40

c* 10 16.8 13.9 1.84 2.2 0.028 0.10 0.022 0.009 0.33

D

* 10 18.6 10.9 1.01 2.2 0.020 0.13 0.002 0.003 0.29

E* 1

1

17.4 13.9 1.58 2.5 0.016 0.17 0.021 0.021 0.48

p

* 1

1

, 1 2 17.2 13.5 1.86 2.3 0.070 0.03 0.024 0.016 0.58

G*13,14 17.3 12.2 1.55 2.2 0.051 0.03 0.026 0.014 0.68

H

* i ° 18.0 10.5 1.10 2.1 0.016 0 . 14-

0.17

0.020 0.015 0.54

I 10 18.0 12.0 1.00 2.1 0.024 0.18 0.003 0.004 0.30

18.0 11.8 1.01 3.0 0.023 0.19 0.002 0.005 0.29

18.2 12.1 1.03 4.1 0.020 0.19 0.002 0.005 0.28

18.5 13.8 1.05 2.0 0.021 0.17 0.002 0.004 0.30

18.2 14.0 1.02 3.0 0.022 0.18 0.002 0.005 0.30

17.9 14.2 1.00 4.0 0.022 0.19 0.002 0.004 0.30

j*10 18.0 10.5 1.10 2.1 0.016 0 . 15-

0.16

0.020 0.015 0.54

Ktio 17.6 13.9 0.97 2.7 0.021 0.18 0.019 0.006 0.54

•fa commercial heat

t electroslag remelted

77
In an earlier study, only the effects of grain size and N content

upon cry were assessed. 3 In accord with the well-known Hall—Petch grainsize effect on yield strength and the higher temperature results of

Norstrom, 15 the form of the equation fitted to the data was

ay

” ao + aN* [ N ] + k'd"*5

. (1)

where [N] is N content (wt.%), d is the grain size (/im) , and aQis the

yield strength in the absence of N with infinite grain size. The contribution from C is expected to be about one-half that of N, 5 but since it

is best to hold C content low (-0.03 wt.%) to minimize sensitization

during welding, the alloys used in this study did not have a sufficient

range of C content to determine this coefficient accurately. In equation (1), the [C] dependence is included in the constant term, aQ

.

To determine whether the addition of a term expressing the dependence upon Mo would lower the SD, the following equation was also fitted

to the data:

ay °o + + aMo -[Mo], (2)

where [Mo] represents Mo content in wt.%. Graphical analysis, especially

of the I alloys (see tables I and II) , had indicated an additional interactive effect of Ni and Mo upon ay.; 16 this was expressed in a trial form:

ay

- aQ + aN

° [N] + k-d^ + aMo -[Mo] + aNr [Ni] + ar [Mo][Ni], (3)

where [Ni] represents Ni content in wt.%. The Ni term was included for

completeness. The units of the coefficients in equations (1) through (3) are as follows: aQ> MPa; ajj, MPa/wt.%; k, MPa^/im^; ajjQ , a^i, MPa/wt.%;

a^, MPa/(wt.%) 2

.

The dependence of Kjc

(J) upon yield strength at 4 K has been observed previously in our laboratory, 1 * 17 and a recent study of the effect

of Ni on toughness of AISI 304- type alloys showed a strong dependence. 8 To assess the relative importance of inclusion spacing with respect to

these two variables, the following equation was previously fitted to a smaller set of data from 316- type alloys: 3 * 4

Kic(J) * + ka

-ay + kfji*[Ni] + k^-A, (4)

where kQ is a constant.

Another expression for the dependence of these parameters that

better expresses the form of current theories 18 is

Kic(J)-V [N1) k°

xl> - where fi is an exponent to be determined in a least- squares fit. The

units of the coefficients in equations (4) and (5) are: kQ , MPa-m^;

ka , m , kfj^, MPa m^/wt . % ; k_\, MPa/mm; kg, (MPa) 2 • 10 3 ^/(wt . %)m^ _ii

.

(5)

78
Both equations (4) and (5) were fitted to the expanded data set. Fig- ure 1 presents a plot of Kj c (J)ay/[Ni] versus A for this data set. The

letter keys to the symbols in this and succeeding figures are given in

table I. For clarity, overlapping data points are omitted from all

figures

.

INCLUSION SPACING, X, mm

Figure 1. The quantity KIc (J) -ay

/[Ni] versus A for 316- type alloys

4 K. The letter keys to the symbols are given in

tables I and II.

79
RESULTS AND DISCUSSION

The result of fitting equation (1) to the set of 99 measurements is:

<7y

(MPa) - 429 + 2500 [N] + 740d" ij

. (6)

The SD of the fit to the data is 45 MPa; the SDs of the three coefficients are 20, 90, and 116, respectively. When we consider their SDs,

these coefficients are in agreement with the coefficients found in

the earlier study for which only 26 measurements were available. Equation (1) with the substitution of [N]^/^ or [N]^/^ for [N] was also

fitted to the data; the result was a substantially higher SD. Following

the work of Norstrom at ambient and higher temperatures, 15 the term

[NJ'd*4 was added to equation (1). This expression also resulted in a substantially higher SD.

Fitting equation (2) to the data gave

oy(MPa) - 316 + 2370 [N] + 790d" H + 53.8[Mo]. (7)

The SD of the fit to the data is 40 MPa; the SDs of the four coefficients

are 28, 84, 102, and 10.2. The constant term, aQ , has decreased because

of the addition of the [Mo] term. The best-determined coefficient (lowest fractional uncertainty) is that of the [N] term; this term also con- tributes the most to ay for 316N-alloy composition. Figure 2, which

presents a plot of (ay — 790d" i

* — 53.8[Mo]) versus [N] , illustrates the

strong dependence of Oy upon [N]

.

Fitting equation (3) to the data set results in a change in the

magnitude of the [Mo] term, as would be expected, but also changes both

the sign and the magnitude of the constant term:

ay(MPa) - -689 + 2510[N] + 823d"*4 + 436 [Mo] + 78.1[Ni] - 30.0[Mo] [Ni] . (8)

The SD of the fit is 34 MPa, and the SDs of the six coefficients are 225,

80, 96, 103, 17.1, and 7.7, respectively. Figure 3 presents the measured

ay values plotted against the values calculated from equation (8)

.

The

straight line in this figure corresponds to agreement between measured

and calculated values, and the scatter band represents two SDs about the

line

.

Equations without the last term of equation (3) or with the first

three terms and an [Ni] term only have larger SDs when fitted to the data

set. However, fitting a subset of the data (I) plus 304-alloy data from

heats supplied by the same manufacturer (alloys H in reference 4) to

equation (3) gave different signs, magnitudes, or both for the constant

term and the last three terms. Most of the Mo variation occurs in subset I of the data, so at present, the correct form of the interaction

terms is considered uncertain. Further analysis 16 of this subset of the

80
(CTy-790d''

/j

-53.8[Mo]),

MPa

Figure 2. Tensile yield strength at 4 K for 316 -type alloys less the

expected contributions from grain size and Mo content [calculated from equation (7)] versus N content. The letter

keys to the symbols are given in tables I and II.

81
MEASURED

TENSILE

YIELD

STRENGTH,

CTy,

MPa

Figure 3. Measured tensile yield strength at 4 K of 316-type alloys

versus tensile yield strength calculated from equation (8)

.

The scatter band represents ± 2 SDs. The letter keys to the

symbols are given in tables I and II.

82
data has shown a saturation of the effect of Mo on cry at 3 wt.%; this

nonlinear effect is not considered in the present linear analysis. The

negative coefficient, a does qualitatively reproduce other results of

reference 16, which show a decrease in ay with increased Ni if Mo is

present. For Mo > 2.6 wt.%, equation ( 8) gives a small decrease in ay

with increased Ni. In general, the effect of Ni on ay is expected to"

be small.

The fit of equation (4) to the data gives:

KIc (J) (MPa-m^) - 150 - 0.314a

y

+ 16.6[Ni] + 2300A, (9)

where the SD of the fit of the data is 30 MPa-m^ and the SDs of the

coefficients are 62, 0.034, 3.5 and 207. When the SDs are considered,

these coefficients again are in good agreement with the results from the

earlier work that was based on only 19 Kj c

(J) measurements. Figure 4

presents the measured Kjc(J) values plotted against the values calculated

from equation (9) ; the straight line corresponds to agreement between

measured and calculated values , and the scatter band represents two SDs

about the line. It is surprising that A correlates so well with Kj c (J)

since it does not represent accurately the volume fraction or shape factor of inclusions (long MnS stringers are equated with small, spherical

inclusions in the simplified counting method). Furthermore, inaccuracies

in A, due to inhomogeneities and counting statistics, may be as large as 30%, particularly in the cleaner alloys.

Equation (5) was also fitted to the data using a linear leastsquares program; the resulting equation is

Kic (J)*y/ [Ni] " (9.56- 10 4 )A°

688

, (10)

where the SD of the predicted Kj c

(J) is 30 MPa-m^, comparable to the

SD of equation (9). The SD of the coefficient of A is 0.077. The form

of this expression may be more easily compared with theory than equation (9), which is simply a linear combination of terms. An equation

similar to (5), Kjc

(J) - k£ • [Ni] a - A^/(ay)T, was also fitted to the data.

The exponents a and 7 were found to be close to unity, and the exponent

P was 0.64, similar to the power of about 2/3 in equation (10). The

difference in repeated measurements of specimens from the same alloy

sample is often 20 to 30 MPa-m^, so SDs below this range may not be

attainable. (See table I for examples.) Also, the effect of sidegrooving during J-integral testing has recently been documented . 19 This

research raises the possibility that earlier NBS measurements not using

side grooves may overestimate the true, plane-strain fracture toughness

for lower strength alloys.

Equations (7) and (9) can be used for 316- type-alloy design. For

example, to achieve a oy

of 1000 MPa with [N] - 0.16 wt. , and [Mo] - 3.0 wt.% (316LN alloy), a grain size of 38 (+29,-14) pm is required.

83
PREDICTED FRACTURE TOUGHNESS, K,C (J), MP<rm'/j

Figure 4. Measured fracture toughness at 4 K of 316- type alloys versus

fracture toughness calculated from equation (9)

.

The scatter

band represents ± 2 SDs. The letter keys to the symbols are given in tables I and II.

84
If [N] - 0.20 wt.% (316LHN alloy), the predicted grain size is 742 (+561,

-263) /im; thus, the large contribution of the [N] term to cr y

outweighs

the grain-size effect.*

To achieve a Kj c

(J) of 200 MPa-m^ with oy

- 1000 MPa would require,

for the 316LN alloy, that A - 0.09 (± 0.05) mm if [Ni] - 10 wt.%, and

A - 0.06 (± 0.03) mm if [Ni] - 14 wt.%, the maximum allowed under the

AISI 316 specifications. The uncertainties are determined in the same manner as above. These inclusion spacings are similar to those found

for the commercial heats in this study (see table I). However, at present, inclusion spacing is not a conventional specification parameter in

commercial alloy production. Further correlation with other parameters

is required.

SUMMARY

Matrices of NBS measurements of ay

and Kj c(J) at 4 K that include

numerical data on alloying, refining, and processing variables were ana- lyzed using linear regression methods. The results of the analysis cor- roborated and extended the results of an earlier study carried out when

fewer data were available. 3 ' 4 The AISI 316- type alloys tested were obtained from a variety of sources and included both commercial and laboratory heats

.

Tensile yield strength : Provisional equations were developed for

estimating oy

at 4 K from [N] , d"*4

, and [Mo] . These equations were derived for a range of [N] from 0.03 to 0.20 wt.%, a range of d from 11

to 242 /im, and a range of [Mo] from 2.0 to 4.0 wt.%. Results from fitting an equation for oy

with an interactive term for [Ni] and [Mo] were

inconclusive.

Fracture toughness : Provisional equations for estimating Kj c (J)

at 4 K from <7y

, [Ni] , and A were determined. In addition, the quantity

Kic(J) -^y/tNi] (fracture toughness normalized for the effects of yield

strength and Ni content) was found to vary approximately as the 2/3 power

of A. These equations were developed for a range of A from 0.04 to

0.11 mm and a range of [Ni] from 10.5 to 14.2 wt.%.

The uncertainties noted in parentheses for these grain sizes are determined by assuming the fractional error in d'

1* is given by the square

root of the sum of the squares of the fractional errors in the terms of

the equation. The fractional error of each term is obtained by dividing

the SD of the coefficient by the value of the coefficient.

85
ACKNOWLEDGMENTS

The research was partially supported by the Office of Fusion Energy,

U.S. Department of Energy, V. Der, Project Monitor, We are indebted to

L. A. Delgado, R. P. Walsh, P. T. Purtscher, and L. L. Scull for assistance in obtaining the data and carrying out the analysis.

REFERENCES

1. R. L. Tobler and R. P. Reed, in: Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures — III, Internal

Report 80—1627, Ed. R. P. Reed (U.S. National Bureau of Standards,

Boulder, Colorado, 1980) p. 15.

2. R. L. Tobler, D. T. Read, and R. P. Reed, in: ASTM STP 743 (American Society for Testing and Materials, Philadelphia, Pennsylvania,

1982), p. 250.

3. N. J. Simon and R. P. Reed, J. Nucl. Mater. 141-143 (1986) 44.

4. N. J. Simon and R. P. Reed, in: Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures — IX, Internal Report

86—3050, Ed. R. P. Reed (U.S. National Bureau of Standards, Boulder,

Colorado, 1986) p. 27.

5. R. P. Reed and N. J. Simon, in: Advances in Cryogenic Engineering —

Materials, vol. 30, Eds. A. F. Clark and R. P. Reed (Plenum Press,

New York, 1984) p. 127.

6. R. P. Reed, N. J. Simon, P. T. Purtscher, and R. L. Tobler, in:

Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures — IX, Internal Report 86—3050, Ed. R. P. Reed (U.S.

National Bureau of Standards, Boulder, Colorado, 1986) p. 15.

7. T. Sakamoto, Y. Nakagama, I. Yamauchi, and T. Zaizen, in: Advances

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

R. P. Reed (Plenum Press, New York, 1984) p. 137.

8. R. P. Reed, P. T. Purtscher, and K. A. Yushchenko, in: Advances in

Cryogenic Engineering — Materials, vol. 32, Eds. R. P. Reed and

A. F. Clark (Plenum Press, New York, 1986) p. 43.

9. N. J. Simon and R. P. Reed, Preliminary Draft, Structural Materials

for Superconducting Magnets, III. AISI 304 Stainless Steel, (U.S.

National Bureau of Standards, Boulder, Colorado, 1985); unpublished.

10.

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

results on AISI 316 -type alloys, U.S. National Bureau of Standards,

Boulder, Colorado, unpublished data (1987).

86
11 . D. T. Read and R. P. Reed, in: Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures — II, Internal Report

79—1609, Eds. F. R. Fickett and R. P. Reed (U.S. National Bureau of

Standards, Boulder, Colorado, 1979) p. 79.

12. R. L. Tobler, R. P. Reed, and D. S. Burkhalter, in: Advances in

Cryogenic Engineering — Materials, vol. 26, Eds. A. F. Clark and

R. P. Reed (Plenum Press, New York, 1980) p. 107.

13. R. L. Tobler, R. P. Mikesell, R. L. Durcholz, C. W. Fowlkes, and

R. P. Reed, in: Semi-Annual Report on Materials Research in Support

of Superconducting Machinery, Internal Report 74—359, (U.S. National

Bureau of Standards, Boulder, Colorado, 1974) p. 182.

14. R. L. Tobler and R. P. Reed, in: Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures — IV, Internal Report

81—1645, Eds. R. P. Reed and N. J. Simon (U.S. National Bureau of

Standards, Boulder, Colorado, 1981) p. 147.

15. L.-A. Norstrom, Met. Sci. 11 (1977) 208.

16. P. T. Purtscher, R. P. Walsh, and R. P. Reed, in: Advances in Cryogenic Engineering - Materials, vol. 34, Eds. A. F. Clark and R. P. Reed (Plenum Press, New York, 1988) p. 191.

17. R. L. Tobler, T. A. Siewert, and H. I. McHenry, in: Materials

Studies for Magnetic Fusion Energy — IX, Internal Report 86—3050,

Ed. R. P. Reed (U.S. National Bureau of Standards, Boulder,

Colorado, 1986) p. 239

18. J. P. Hirth and F. H. Froes, Metall. Trans. 8A (1977) 1165.

19. T. Ogata, K. Ishikawa, T. Yuri, R. L. Tobler, P. T. Purtscher, R. P. Reed, T. Shoji, N. Nakano, and H. Takahashi, in: Advances in Cryogenic Engineering - Materials, vol. 34, Eds. A. F. Clark and R. P. Reed (Plenum Press, New York, 1988) p. 259.

87
'
CLEAVAGE -LIKE FRACTURE ALONG SLIP PLANES

IN Fe—18Cr—3Ni—13Mrt—0 . 37N AUSTENITIC STAINLESS STEEL

AT LIQUID -HELIUM TEMPERATURE*

R. L. Tobler

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

and

D, Meyn

Materials Science and Technology Division

Naval Research Laboratory

Washington, D.C.

Contrary to usual expectations for the face-centered cubic

crystal structure, some nitrogen- alloyed austenitic stainless

steels are brittle at temperatures approaching absolute zero. In this communication, results of plane-strain fracture tough- ness and fractographic tests are presented for an Fe—18Cr—3Ni—

13Mn-0.37N steel tested at liquid-helium temperature (4 K)

.

The strength of this steel increases significantly at cryogenic

temperatures owing to high-nitrogen content. At 4 K, the. yield

strength reaches 1540 MPa. Low toughness (Kj c - 70 MPa -in’

4

) and linear-elastic behavior is observed in compact specimens

at this temperature, in association with a brittle failure

mechanism consisting of transgranular facets. Valid fracturetoughness (Kj c ) measurements at 4 K are reported, the cleavage - like failure mechanism is described, and the crystal plane indices of the transgranular facets produced during fatigue and

fracture are identified by x-ray diffraction and slip- line

traces

.

* Submitted to Metallurgical Transactions

.

89
The test material was a 25 . 4-mm- thick plate of Fe—18Cr—3Ni—13Mn—

0.37N austenitic stainless steel in the commercially annealed condition:

1080°C for 1 h and water quenched. The mill chemical analysis in weight

percent is : Fe-18 . 09Cr-3 . 26Ni-13 . 22Mn-0 . 37N-0 . 038C-0 . 12Mo-0 . 52Si-0 . 005S0.028P. The material in the as -received condition had an average aus- tenite grain diameter of 48 /xm and a microstructure free of carbide

precipitation. The mechanical properties are listed in Table 1.

Table 1. Mechanical properties of as-received Fe—18Cr—3Ni—13Mn—0 . 37N

austenitic steel [1].

Temperature

(K)

Yield

Stress

(MPa)

Ultimate

Strength

(MPa)

Elongation

(%)

Reduction

of Area

(%)

Rockwell

Hardness

(Rb)

295 400 796 56 53 93

4 1540 1811 4 26 -

Toughness tests according to Methods E 399-83 [2] were performed

using compact specimens 25. 4-mm thick (B) and 50 . 8 -mm wide (W)

.

Two

specimens were machined from the as -received material, and one from

laboratory- annealed material. The laboratory anneal (1093°C, 37 h, water

quenched) was applied to a sample of this steel to enlarge its grains.

All specimens were fatigue -cracked at stress -intensity factor ranges from

20 to 50 MPa-m1* and then fractured at 4 K, The average grain diameter

after the laboratory anneal was large enough (about 1 mm) to enable Laue

x-ray analysis of individual crystal facets situated on the fracture sur- face of the polycrystalline aggregate.

The load-versus- deflection curve for the Fe—18Cr-3Ni-13Mn—0 . 37N

specimen in the as-received condition at 4 K is shown in Fig. 1. A

type-I [2] linear-elastic behavior was observed, with only slight non- linearity prior to fast fracture. The typical fracture -specimen surface

(Fig. 2) displayed no shear lips. Both the fatigue and overload failure

zones were both covered by smooth, highly reflective facets producing a macroscopic appearance; such facets are commonly seen in body-centered

cubic (bcc) metals fractured at low temperatures. There was a relatively

large degree of fatigue crack curvature: the edge crack lengths were

about 14 percent shorter than the three -point average crack lengths calculated according to Methods E 399-83. Thus, the crack front curvature

was sizable, but it did not exceed the 15 percent limitation set by sec- tion 8.2.2 of Methods E 399-83 [2]. Magnetometer and x-ray examination

of the 4-K fracture surface revealed no detectable bcc martensitic phase;

the sensitivities were 0.2 and 3 percent, respectively.

90
Figure 1. Load-versus- deflection test record at 4 K for a chevron

notched specimen of Fe—18Cr—3Ni—13Mn—0 . 37N steel in the

as-received condition.

The fracture -toughness (Kj c ) measurements for the as -received ma- terial are listed in Table 2. Also listed are specimen dimensions and

test-validity criteria, as defined in Methods E 399-83. Although the

fracture toughness of the Fe—18Cr-3Ni—13Mn—0 . 37N steel seems low for an austenitic alloy, in fact it meets expectations for steels of this

strength level, based on the known inverse trend of Kj c versus Oy [1]

.

The results for a chevron-notched and a straight- through-notched specimen

proved indistinguishable: Their Kj c values were nearly equal, and their

crack-front curvatures were identical. Thus, the crack-starter notch

type had no influence on the final crack- front shape.

91
Figure 2.

Macroscopic fracture sur- face of chevron-notched

Fe—1 8Cr—3Ni—1 3Mn—0 . 37N

steel after testing:

A) 76 -K fatigue zone,

B) 4-K fatigue zone,

C) 4-K overload failure zone,

D) final fracture zone at

T > 4 K.

10 mm

92
Table 2. Fracture - toughness data for as-received Fe—18Cr-3Ni-13Mn-0 . 37N

austenitic steel at 4 K.

Speci- men

Number

Specimen

Thickness

,

5

(mm)

Specimen

Width,

W

(mm)

Average

Crack

Length , a

(mm)

Relative

Crack

Length

,

a/W

Load

Ratio

^max/^Q

Required

Thickness

2 . 5 (KIc/

o

y

) 2 (mm)

Fracture

Toughness

,

Klc , (MPa-m4

)

1* 25.400 50.725 25.975 0.512 1.05 5.2 70

2t 25.425 50.600 26.650 0.527 1.01 5.3 71 x chevron-notch design

t straight- through-notch design

Our Kj c values agree well with the average Kq value (71 MPa-m^)

reported by Read and Reed [1] for the same plate of steel. Their Kq

values are reported as invalid because they did not satisfy the crackfront -curvature requirements of Methods E 399-74, the current standard at

the time. Although similar curvatures occurred in the present study, our measurements are valid because the curvature tolerances were relaxed in

the updated Methods E 399-83 standard.

Figures 3 and 4 show facets produced in the coarse-grained specimen

of Fe—18Cr—3Ni—13Mn—0 . 37N steel that was fatigued and fractured at 4 K.

The facet planes in both zones were oriented at angles from 90° to 53°

with respect to the loading axis, or 0° to 37° from the macroscopic plane

of fracture. As shown at higher magnification in Figs. 5 and 6, the

facets displayed steps and slip-line markings, but no river patterns,

such as those observed on classical cleavage facets for ferritic steels.

Three facets in the fatigue- failure zone of the coarse-grained

specimen were examined by Laue x-ray diffraction, and all three were

identified as {111} planes in an fee structure. The crystallographic

directions in these {111} planes were apparently randomly oriented with

respect to the crack propagation direction. The fee structure underlying

the facets was verified by careful structure factor analysis of the Laue

back-reflection patterns and comparison with known fee and bcc patterns.

Overload-zone facets for the coarse-grained specimen were also identified as {111} planes on the basis of the 60° angle between the slipline traces and between facet edges. Again, three facets in this zone were x-rayed. The planar indices could not be confirmed by x-ray diffraction because the diffraction technique produced no spot pattern, only

general blackening due to microplastic deformation induced at fracture.

Consequently, the crystallographic directions for overload facets could

not be determined.

93
0.24 mm

0 .1 mm

Figure 3. Scanning -electron microscope photographs of crystallographic

facets on the fatigue failure zone of coarse-grained Fe—18Cr—

3Ni—13Mn—0 . 37N steel fractured at 4 K: A) facet no. 2 , B) facet no. 1.

94
0.1 mm

B

0.2 mm

Figure 4. Scanning-electron microscope photographs of crystallographic

facets on the overload fracture zone of coarse-grained Fe18Cr—3Ni—13Mn-0 . 37N steel fractured at 4 K: A) facet no. 4,

B) facet no . 5

.

95
AO 4111

2 pm

Figure 5. Scanning-electron microscope photographs of microscopic steps

and fine slip- line traces on a fatigue-zone facet (facet no. 2 of Fig. 3): A) smooth area, B) same facet, but showing slip

markings in detail.

96
2 pm

Figure 6. Scanning- electron microscope photographs of slip- line traces on an overload-zone facet (facet no. 4 of Fig. 4): A) plane

area with intersecting slip lines, B) area with "striations"

caused by regularly spaced slip bands.

97
According to a simple model, the occurrence of ductile or brittle

fracture at a given test temperature depends on whether the yield

strength (cry) or the cleavage strength (ac ) is reached first [3,4]. For

metals, ac is relatively high and temperature insensitive. Critical

factors, then, are the magnitude of Oy and its rate of increase at low

temperatures

.

For most bcc metals and alloys, including ferritic steels, Oy is

strongly temperature dependent. As the test temperature is lowered, the

lattice resistance to slip increases until, at some temperature, ct c is

reached before slip can occur. As a result, cleavage is common in bcc

materials at low temperatures. There are exceptions to the rule, however. For example, zone-refined, high-purity bcc iron remains ductile in

tensile tests at temperatures as low as 4 K [5].

In the case of most fee metals and alloys, including austenitic

stainless steels, Oy is low and not strongly temperature dependent.

Early observations led to the traditional view that "fee metals do not

fracture by cleavage” [3], and we overgeneralized that "slip occurs readily in fee metals at all temperatures, their shear strengths are always lower than their cleavage strengths, and they always fail in a ductile manner" [4]

.

In reproof, a number of exceptions have not appeared as our experience has grown [6—19]. For example, rhodium and

iridium cleave in inert environments at room temperature [6,7]; nickel

and AlTM6Zn^3Mg cleave in liquid metals at various temperatures [8]

;

and

strong nitrogen- alloyed austenitic stainless steels cleave at liquidnitrogen temperature [9]

.

Regarding failure mechanisms, tensile loading of iridium in inert

environments at room temperature, or of aluminum in liquid mercury, produces cleavage on {001} planes [6,8]. Corrosion- fatigue cracking in aluminum alloys also produces cleavage on {001} planes [11]. Cleavage on

{111} planes has been observed in fee metals under conditions of fatigue

in alloys with coarse-grained structures and limited cross-slip [12].

Also, cleavage- like slip-plane cracking on {111} proves to be the usual

crack- growth mechanism for nickel -base superalloys fatigued at room

temperature [ 13—15 ]

.

For austenitic steels, Caskey [9] reports cleavage for an Fe—

17.44Cr—15.31Mn-(0.4—0.6)N alloy in tensile tests at 78 K. The fracture

surface displayed cleavage facets lying along coherent twin boundaries,

transverse to the tensile axis. The fracture progressed by brittle

microcracking across individual grains, followed by ductile fracture of

the intervening material. The faceted failure mechanism was attributed

to deformation faulting and to e -phase formation. Others reported similar brittleness in Fe—Cr—Mn—Ni steels at low temperatures [16—19] ; this was attributed to several possible mechanisms, including the formation of

brittle martensites and the high localization of strain arising from

deformation faulting.

The Fe—18Cr—3Ni—13Mn—0 . 37N steel of this study is related to the

Fe—Cr—Mn—Ni steels discussed above. They were all developed for the

98
economical replacement of nickel by manganese and nitrogen. Manganese

increases the solubility for nitrogen, the most potent austenitestrengthening element at low temperatures. Thus, the steel of this

study contains 0.37% nitrogen, and its yield strength more than triples

between 295 and 4 K, promoting brittle fracture, as mentioned above.

Various definitions of cleavage have been offered, but ASTM Standard E 616-82 on terminology defines cleavage (strictly, crystallographic

cleavage) as "separation of a crystal along a plane of fixed orientation

relative to the three-dimensional crystal structure within which the

separation process occurs, with the separation process causing the newly

formed surfaces to move away from one another in directions containing

major components of motion perpendicular to the fixed plane" [20] . The

facets on our Fe-18Cr—3Ni-13Mn-0 . 37N steel surfaces satisfy this definition of cleavage, even though the failure mechanism is evidently slipband cracking.

Fractographically , slip band-cracking is similar to tensile cleavage, but the distinction is noteworthy. If the brittle fracture mechanism for our steel were pure tensile cleavage, then the cleavage plane

would be {001}, not {111}. The usual tensile -cleavage plane in fee

alloys is {001}, as discussed above, whereas {111} cleavage has been

observed only tinder conditions giving rise to slip -band cracking, such as that which occurs during Stage-I fatigue-crack propagation [13-15]. Such

observations, which are common for high- temperature nickel alloys, indicate features strongly resembling those of the present study, including

cleavage planes at high angles to the stress axis. Slip-band or shearband cracking has been identified as the mechanism for {111} cleavage in

the nickel-alloy studies, and it is distinct from simple slipping off, or "slip-plane decohesion," which causes chisel-edge fractures in single

crystals

.

In general, slip-band cracking is facilitated by planar slip (low

stacking- fault energies, short-range order, and shearable precipitates)

and by coarse grain size. The sequence of events for si ip -band cracking

in our steel might be: a narrow, well-defined shear (slip) band forms in a suitably oriented grain, and because of the difficulty of activating

out-of-plane secondary slip and the restriction of shear to the original

plane by metallurgical factors, shear strain continues to accumulate in

the narrow band. The accumulation of dislocation- induced defects in this

band weakens it to the extent that the component of tensile force on the

band causes final separation in the tensile mode, and this gives rise to

the observed cleavage- like appearance.

Given the complex character of the stress and strain states and the

various grain orientations in the crack- tip region, it is not surprising

that significant amounts of shear strain might accumulate in a few grain:;

on slip bands oriented nearly perpendicular to the gross-stress axis.

The amount of shear-strain accumulation needed to facilitate tensile

cracking on such a band would be a function of its orientation, which

determines the magnitude of the resolved tensile component. The overall

slip-band cracking mechanism, seen in this light, becomes an interplay

99
tensile component: the more nearly perpendicular the slip band is to the

gross-tensile-load axis, the higher the local tensile stress across the

slip band and the less shear-strain accumulation needed to trigger separation. One would expect that some range of angles between 90° and 45°

(with respect to the gross-stress axis) would be most favorable for slip

band cracking, as defined here.

Finally, Koss and Chan [21] have analyzed the conditions for crack

advance along planar slip bands in non-work-hardening materials. They

assume a condition in which slip is allowed to occur in the plane of the

crack, with plastic displacements restricted to shear that is parallel to

the crack plane. Then, once a crack forms in a coplanar slip band, it

follows that it is difficult to activate secondary slip with shear displacements inclined to the crack plane. The result: a large local shear

strain, combined with large normal or hydrostatic stresses near the crack

tip, a condition conducive to easy crack propagation and low toughness.

This analysis also shows that even if cross-slip could occur, it would

not be able to relax the large normal stress on slip bands coplanar with

the crack [22]. Thus, the stress state that exists under the assumed

conditions is the paramount factor, and the role of specific dislocation

arrangements or the weakening of slip bands by accumulated shear is deemphasized. This model helps to rationalize some features of slip-band

cracking in nickel-, aluminum-, and titanium-based alloys [21]. Similarly, it gains relevance to the austenitic steel of our study in view

of its decreased work-hardening capability, increased strength, and increased planarity of slip at 4 K.

ACKNOWLEDGMENT

This study was supported by the U.S. Department of Energy, Office of

Fusion Energy.

REFERENCES

1. D. T. Read and R. P. Reed: The Metal Science of Stainless Steels,

E. W. Ceilings and H. W. King, eds., Metallurgical Society of the

AIME, New York, 1979, pp. 92-121.

2. Standard Test Method for Plane-Strain Fracture Toughness of Metallic

Materials, ASTM Designation E 399-83, 1987 Annual Book of ASTM Standards, Section 3, American Society for Testing and Materials, Philadelphia, 1987, pp. 680-715.

3. A. S. Tetelman and A. J. McEvily, Jr.: Fracture of Structural Materials, Wiley, New York, 1967, p. 212.

4. D. A. Wigley: Mechanical Properties of Materials at Low Temperatures, Plenum, New York, 1971, pp. 94—95.

5. Z. S. Basinski and A. Sleeswyk: Acta Metall., 1957, vol. 5,

pp. 176-179.

6. S. S. Hecker, D. L. Rohr, and D. F. Stein: Metall. Trans. A,

vol. 9A, 1978, 481-487.

100
7. C. Gandhi and M. F. Ashby: Scr. Metall., 1979, vol. 13, pp. 371—

379.

8. C. F. Old, Met. Sci. , 1980, vol. 14, pp. 433-440.

9. G. R. Caskey, Jr.: Fracture of Fe—Cr-Mn austenitic steel, Report

DP-MS-78-68, E. I. du Pont de Nemours, Aiken, South Carolina, 1978,

pp. 1-21.

10. S. P. Lynch: Met. Sci., 1981, vol. 15, pp. 463—467.

11. G. G. Garrett and J. F. Knott: Acta Metall., 1975, vol. 23,

pp. 841—848.

12. K. R. L. Thompson and J. V. Craig: Metall. Trans. A, 1970, vol. 1A,

pp. 1047-1049.

13. D. J. Duquette, M. Gell, and J. W. Piteo: Met. Trans. A, 1970,

vol. 1A, pp. 3107-3115.

14. D. J. Duquette and M. Gell: Met. Trans. A, 1971, vol. 2A, pp. 1325-

1331.

15. M. Gell and G. R. Leverant: Acta Metall., 1968, vol. 16, pp. 553—

561.

16. F. W. Schaller and V. F. Zackay: Trans. ASM, 1951, vol. 51,

pp. 609-628.

17. C. M. Hsiao and E. J. Dulis: Trans ASM, 1960, vol. 52, pp. 855—877.

18. J. C. Shyne, F. W. Schaller, and V. F. Zackay: Trans. ASM, 1960,

vol. 52, pp. 848—854.

19. J. D. Defilippi, K. G. Brickner, and E. M. Gilbert: Trans. Metall.

Soc. AIME, 1969, vol. 245, pp. 2141-2148.

20. Standard Terminology Relating to Fracture Testing, ASTM Designation

E 616-82, 1987 Annual Book of ASTM Standards, Section 3, Vol. 03.01,

American Society for Testing and Materials, Philadelphia, 1987,

p. 864.

21. D. A. Koss and K.

1252.

S. Chan: Acta Metall .

,

1980, vol. 28, pp. 1245-

22. K. S. Chan: Acta Metall .

,

1987, vol. 35, pp. 981-987.

101

FRACTURE BEHAVIOR OF 316LN ALLOY

IN UNIAXIAL TENSION AT CRYOGENIC TEMPERATURES*

P. T. Purtscher, R. P. Walsh, and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The fracture behavior of an austenitic stainless steel,

commercial -grade 316LN alloy, at cryogenic temperatures was studied by careful examination of the fracture surfaces and

polished cross sections through the fracture surfaces of round

tensile specimens. The fracture was ductile (a dimpled rupture

process) and was controlled by inclusions (MnS stringers and

smaller spherical particles) in the structure. The main effect

of the MnS stringers was to decrease the percent reduction in

area in the tensile test at 4 K. The true stress and strain

at fracture varied as a function of test temperature; they were

related to the nucleation of microvoids around the spherical

particles

.

INTRODUCTION

Many studies 1 " 4 have characterized the mechanical properties of aus- tenitic stainless steels at 4 K. Typically, the previous studies looked

at the toughness, tensile yield strength, and serrated yielding behavior,

but paid little attention to the tensile fracture behavior. The tensile

fracture behavior can provide additional information on the deformation

and fracture processes and increase our understanding of fracture in more complex stress states.

In this study, we examined the fracture behavior of a commercial

grade 316LN austenitic stainless steel at cryogenic temperatures.

Specifically, we considered the role of inclusions and the effect of

*In Advances in Cryogenic Engineering — Materials

,

vol. 34, Plenum,

New York, 1988, pp. 379-386.

103
temperature. Round tensile specimens with a uniform gage length fractured in uniaxial tension were chosen for this initial study to simplify

the interpretation of the results. The fracture surfaces were examined

carefully and related to the features found in longitudinal cross sec- tions through the fracture surfaces. The mechanical properties of the

commercial alloy at 4 K are compared with the same properties from a laboratory heat of 316LN with a different inclusion content.

MATERIALS AND PROCEDURE

A commercial -grade 316LN alloy was supplied as 50-mm- thick plate,

mill annealed at a minimum temperature of 1038°C. The 4-K mechanical

properties of this alloy are reported elsewhere. 8 The tensile properties at 4 and 76 K from NBS tests are listed along with the steel mill's

295-K results in Table I. The 4-K mechanical properties of a laboratory

heat of 316LN are also listed for comparison.

The composition, grain size, and inclusion spacing for the two 316LN

alloys are shown in Table II. The inclusion spacing was determined by

counting all inclusions greater than 0.5 /xm in diameter in a given area,

normalizing that number to a 1 mm2 area, and taking the inverse square

root.

Table I. Tensile Properties of 316LN Alloys

Source

Temperature

<K)

Strength

Yield

(MPa)

Ultimate

Ductility

El.

(%)

R. A.

Commercial* 295 275 581 60 68

Commercialt 76 757 to 761 1430 53 65

Commercial^ 4 845 to 860 1630 to 1640 38 to 46 37 to 48

Laboratory^ 4 840 to 1060 1620 to 1730 36 to 43 49-58

< results from steel mill

t one specimen pulled to failure; four more interrupted tests

$ five specimens pulled to failure

Table II. Composition (wt.%) of 316LN Alloys

C Mn P S Si Cr Ni Mo N

Commercial 0.016 1.1 0.020 0.015 0.54 18.0 10.5 2.12 0.15

heat

Laboratory 0.020 1.0 0.002 0.003 0.29 18.6 10.9 2.2 0.12

heat

104
The commercial alloy contained two distinct types of particles in

the austenitic matrix. Figure 1 shows both types in the etched micro

structure. The first type is a large, elongated MnS inclusion that

runs parallel to the rolling direction. The second type of particle is

spherical and much smaller in size (1 to 2 pm in diameter) . X-ray analysis with an energy -dispersive analyzer showed that these smaller particles contain Al, Ti, and Ca, but we could not definitely identify the

particles as carbide, nitride, oxide, or silicate types . The laboratory

heat did not contain large elongated MnS stringers. All the inclusions

were spherical, approximately 1 to 2 /xm in diameter; all contained Mn, S,

Si, Ca, and Al.

Standard round tensile specimens, 6.2 mm in diameter and 38 mm in

gage length, were machined so that the tensile axis was transverse to the

rolling direction. The tests were conducted in displacement control with

a strain rate of 2 x 10" 4 per second. At 76 K, the strain rate was re- duced one order of magnitude after the maximum load. A total of five

tests were performed in liquid nitrogen (76 K) . Only one specimen was actually broken; the others were stopped at different points past the

Figure 1. Etched microstructure showing the two types of particles:

A — MnS stringers; B - smaller, spherical particles.

105
maximum load. At 4 K, the results of five tests, in which all specimens

were pulled to failure, are reported. The initial strain was determined

from clip-on extensometers . At the end of the test, the strain was determined from the diameter of the specimen and

e - 2 In D0/Df , (1)

where D0 and Df are the original and final diameters, respectively.

RESULTS

Figure 2 shows a schematic of the load-vs . - time output from the

tests. The most obvious difference between the test results is the ser- rated flow of the curve at 4 K. Extensive studies of serrated yielding

have shown that the load drops are associated with local temperature

rises and avalanches of dislocation motion in the region of the temperature rise. 4 ' 6 * 7 The plastic deformation occurs in discrete steps along

the gage length similar to Luders bands.

Figure 2. Schematic curves representing the load-vs . -time behavior of

tensile specimens at 4 and 76 K, where a

f

and 6f are the true

stress and strain to fracture.

106
In general, the yield and flow strengths at 76 K are proportionately

lower than at 4 K. Excepting the serrated yielding, the load-vs . - time curves at the two temperatures have similar shapes until maximum load.

Both curves initially show low work hardening owing to the strain- induced

martensitic transformation. At the maximum load, the specimens at 4 K

fractured with a total elongation of about 43%. At 76 K, the total elongation was slightly higher, and the load decreased about 20% with further

straining before the specimen fractured. Approximately the same amount

of strain- induced martensite formed at each test temperature.

At 76 K, four tests were stopped before fracture so that the progression of events that led to the final fracture could be studied. The

maximum true stress and true strain for each of these specimens is shown

in Figure 2 at approximately the point at which the tests were stopped.

At 4 K, the maximum load could not be detected prior to fracture. The

true stress and true strain for the two points marked on the 4-K curve in

Figure 2 were taken from a representative specimen. The true stress and

strain at fracture were calculated from the final load and final cross- sectional area. The lower values of true stress and strain at 4 K were

calculated from the load at the point on the curve just prior to the

final serration and the cross-sectional area outside the localized neck

in a region of relatively uniform strain.

The fracture surfaces of the broken tensile specimens were examined

in a scanning electron microscope to determine the fracture morphology.

The tensile fracture surfaces of the 76- and 4-K specimens were very

similar. The surface had an outer ring of smooth shear lip and an inner

region with a rougher, "woody" texture. In the higher magnification of

the rough- textured area in Figure 3, large, deep delaminations can be

seen; they varied in size from less than 1 to 2 or 3 ^ in diameter.

The larger dimples had spherical particles in their center that served as the nucleation site. Particles were not observed in the smaller dimples.

However, it is possible that particles less than 0.5 pm in diameter could

exist in the austenitic matrix and serve as a nucleation site for dimple

formation. Some of the small dimples lie on a plane at a 45-degree angle

to the tensile axis; they are either closed or open-ended elongated

dimples. The other small dimples appear equiaxed and are approximately

perpendicular to the tensile axis.

Selected specimens were sectioned longitudinally and mechanically

polished to observe the profile of the fracture surface and the area below the surface. At 4 K in the region outside of the neck, MnS stringers were cracked, but no apparent void growth was observed. In the

necked region (Figure 4) , the MnS stringers produced voids that formed

de laminations . These delaminations are linked by microvoids on the fracture surface. There is no apparent nucleation of the smaller dimples

below the fracture surface. The linking of the delaminations by the

nucleation and the growth of the smaller dimples appears to be a sudden

catastrophic event.

107
50 pm u 10 (jin

Figure 3. High-magnification views of tensile fracture surface

specimen tested at 4 K.

Specimens tested at 76 K had a very similar appearance. The large

MnS stringers had fractured by the time the load reached a maximum, but

there was no apparent void growth. As the load dropped from the maximum

point, delaminations formed around the stringers. The smaller voids that

linked the delaminations on the fracture surface did not appear on the

polished cross sections of the interrupted tests. Presumably, the nu- cleation and growth of the microvoids by tensile loading created a transverse crack in the center of the necked region. The outer ligament then

failed in shear. These presumed events must have occurred between points

1 and 2 at 4 K and points 6 and 7 at 76 K in Figure 2.

DISCUSSION

The results show that there was very little change in the fracture

behavior of the 316LN specimens tested in tension at 76 and 4 K. The

same physical processes occurred at both temperatures. The serrated

yielding did not change the process; it only made the process discontinuous and localized.

We interpreted our observations of the tensile fracture behavior in

terms of two distinct steps: (1) the nucleation and limited growth of

delaminations around the large MnS stringers and (2) the nucleation of a second set of smaller dimples in the ligaments between the large delamination, which occurred suddenly and led directly to the fracture without

any additional applied strain.

108
b

20 um

Figure 4. Magnified view of the longitudinal cross section

from a fracture surface of a 4-K tensile specimen.

109
The large MnS stringers produced the de laminations that effectively

transformed the necked region of the tensile specimen into many smaller

tensile specimens (step 1). If the MnS stringers could be eliminated,

the tensile properties would improve. The magnitude of the improvement

we would expect in the 4-K properties can be obtained by comparing the

properties of this commercial -grade steel with those of the laboratoryproduced heat of 316LN (Table I) . The compositions of the two heats were

similar, except the S and P content of the laboratory heat was approximately one order of magnitude lower. The lower impurity content reflected its larger inclusion spacing. In addition, no MnS stringers were

observed in the laboratory-produced heat. The only mechanical property

that appears to have been affected by the presence of the large stringers

is the reduction in area. The commercial heat with the large stringers

had 20% less reduction in area than the laboratory heat.

Modem steel -making methods have made it possible to produce aus- tenitic stainless steels with smaller and fewer MnS stringers. Our

results indicate that these methods will not significantly improve the

tensile ductilities of the alloy. The additional cost of producing

steels free from large MnS stringers may be justified on the basis of

other test results, such as toughness.

The second step had two components. The small dimples on a plane

that was at a 45-degree angle to the tensile axis can be described as void- sheet coalescence or shear lip. High shear stresses near the free

surface of the delamination were responsible for their formation. The

rest of the smaller dimples formed in step 2 lay on a plane roughly perpendicular to the tensile axis and were produced by the normal stresses

acting along the tensile axis.

The second step was controlled by the formation of the smaller

dimples. The macroscopic true stress and strain present at the point

of fracture varied with the test temperature, as shown in Figure 2:

2473 MPa and 0.52 at 4 K; 3051 MPa and 1.04 at 76 K, To understand the

nucleation of the small dimples (step 2) , we present a simple model to

describe nucleation, similar to one proposed by LeRoy et al. 8 The model

assumes that nucleation occurs when the interfacial stress (ac ) between

the particle and matrix reaches a critical value. The critical stress

in the model is composed of a local stress (o^) , a function of the nu- cleation strain (e n), and the macroscopic stress (am)

ac " ai + am /*N CM V-/

(3)

The critical stress for a tightly bound particle (in two ferritic

steels and a Cu alloy) was determined by Argon and Im9 to be equal to the

Young's modulus (E) of the matrix divided by 120 when the macroscopic

stress is assumed to be equal to the hydrostatic component of the applied

stress. The cry parameter is related to the stress associated with the

dislocations piled up at obstacles, namely the particles in the matrix.

110
Since cross -slip is the mechanism likely to relieve the stress concentration at the front of a pileup, the parameters that affect cross -slip,

such as temperature and stacking -fault energy, are likely to affect

The stresses and strains at fracture (the applied stresses and

strains required for nucleation of the smaller dimples) can be substituted into Equations 2 and 3 in place of the macroscopic stress and the

nucleation strain. Because we are using the total applied stress for the

macroscopic component in Equation 2, then the critical stress we use should be approximately 3 times greater than Argon and Im9 reported, E/40

rather than E/120. If ac does not change with temperature, then the an value can be calculated as a function of temperature, 3500 MPa at 4 K and

1875 MPa at 76 K.

The model predicts that ctjj will decrease as the test temperature is

raised from 4 to 76 K, indicating, as expected, that the stress associated with dislocation pileups is less at 76 K than at 4 K. The nucleation model provides a simple method of relating the macroscopic stresses

and strains to the microscopic process. The parameters ac and a# must be

better defined and correlated with other results before they can be used

reliab ly

.

CONCLUSIONS

1. The fracture behavior of the commercial -grade 316LN alloy at cryogenic^ temperatures was influenced by both the large MnS stringers

and the small spherical particles.

2. The presence of large stringers in this alloy affects the measured

mechanical properties mainly by reducing the percent reduction in

area during the tensile test. The other mechanical properties, percent elongation, yield strength, ultimate strength, and fracture

toughness, appear to be unaffected by the large stringers.

3. The serrated yielding that occurs in this alloy at 4 K and not at

76 K has no discernable effect on the fracture behavior.

4. The fracture conditions, true stress and strain, increase as the test

temperature increases and appear to be controlled by the nucleation

of microvoids at the small, spherical particles. A simple model describes the fracture conditions to be a function of (a) the strength

of the particle-matrix interface and (b) the stress associated with

the dislocations that are piled up against the interface.

ACKNOWLEDGMENTS

This work was partially supported by the Department of Energy, Office

of Fusion Energy. The authors thank Luz Delgado for her assistance in

the mechanical testing and Lawrence Livermore Laboratory for supplying

the commercial -grade alloy.

Ill
REFERENCES

1. R. P. Reed, P. T. Purtscher, and K. A. Yushchenko, in: Advances in

Cryogenic Engineering — Materials , vol. 32, R. P. Reed and A. F. Clark, eds., Plenum Press, New York, 1986, p. 43.

2. K. Suemune, K. Sugino, H. Masumoto, H. Nakajima, and S. Shimamoto,

in: Advances in Cryogenic Engineering — Materials , vol. 32, R. P. Reed and A. F. Clark, eds., Plenum Press, New York, 1986, p. 51.

3. S. Yamamoto, N. Yamagami, and C. Ouchi, in: Advances in Cryogenic

Engineering — Materials , vol. 32, R. P. Reed and A. F. Clark, eds.,

Plenum Press, New York, 1986, p. 57.

4. D. T. Read and R. P. Reed, in: Advances in Cryogenic Engineering —

Materials, vol. 26, A. F. Clark and R. P. Reed, eds., Plenum Press,

New York, 1980, p. 91.

5. R. P. Reed, N. J. Simon, P. T. Purtscher, and R. L. Tobler, in:

Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures — IX . R. P. Reed, ed. , National Bureau of Standards,

Boulder, Colorado, 1986, p. 15.

6. K. Shibata and T. Fujita, Trans. ISIJ . vol. 26, 1986, p. 1065.

7. Z. S. Basinski, Proc. Rov. Soc . . vol. A240, 1957, p. 229.

8. G. LeRoy, J. D. Embry, G. Edwards, and M. F. Ashby, Acta Metall .. vol. 29, 1981, p. 1509.

9. A. S. Argon and J. Im, Metall. Trans ., vol. 6A, 1975, p. 839.

112
EFFECT OF VOID NUCLEATION

ON FRACTURE TOUGHNESS OF HIGH-STRENGTH AUSTENITIC STEELS*

P. T. Purtscher, R. P. Reed, and D. T. Read

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The fracture of seven austenitic stainless steels with

varying Ni and N contents were studied at 4 K. Smooth, 6-mmdiameter tensile specimens and 22-mm-thick compact specimens

were used. Nitrogen content controlled the yield strength and

influenced fracture toughness by its effect on yield strength.

At a constant N content, increasing the Ni content increased

the fracture toughness. Observations of the fracture surfaces

and polished cross sections through the fracture surfaces of

test specimens showed that nucleation controlled the dimpled

rupture fracture process. A critical stress criterion for

nucleation that depends on both the applied stress and strain

was developed and applied to the fracture toughness test. This

fracture criterion explained the increase in fracture toughness

with increasing Ni and the decrease in fracture toughness with

increasing yield strength for strengths over 600 MPa.

INTRODUCTION

The fracture toughness characterizes the initiation of the ductile

fracture process ahead of a fatigue precrack. Two measures of fracture

toughness, Jj c and Kj c , have been used for applications where high

strength and high reliability are required, such as aircraft parts and

pressure vessels. Nevertheless, the basic mechanisms that relate the

microstructure of a material to its fracture mechanics properties are not understood. It is extremely difficult to predict Jj c from the microstructure and strength of an alloy. 1

Austenitic stainless steels are used in low- temperature structural

applications where Jj c is an important design consideration: liquefied

gas storage tanks and supports for superconducting magnets. 2 At cryogenic temperatures, these alloys can have high yield strengths and retair.

a relatively high Jj c

.

* Submitted to American Society for Testing and Materials.

113
Figure 1. Fracture toughness vs. yield strength for structural alloys.

Data for Al, Ti, new steels, and old steels are from Reference 7. Austenitic stainless steel data are from Reference 8.

114
The relationship between strength and fracture toughness of structural alloys has been studied extensively. 3

" 5 Wanhill 6 has summarized

the results of numerous studies on the most common high-strength alloys.

To Vanhill's data, we added the strength -vs .- toughness data for a con- trolled series of austenitic stainless steels whose strength was raised

systematically by increasing the concentration of interstitial alloying

elements 7 (see Figure 1). For these austenitic steels, the two properties are related empirically by the following equation:

KIc (J) - 500 - 0.3 Oy (1)

where Ki c(J) is in MPa-m^; ay, the yield stress, is in MPa. This equation predicts the fracture toughness within 20 MPa- in’*. All the curves

in Figure 1 have similar negative slopes, reflecting the significant

influence of strength on the fracture toughness of high-strength, structural alloys.

Steels fracture in a ductile manner by a dimpled rupture process in

which voids nucleate at discontinuities in the micro s true ture , typically

nonmetallic inclusions, and grow under the influence of the local stress

state. In the uniaxial tensile test, the main direction of void growth

is along the tensile axis until a local neck is formed. The neck introduces triaxial stresses that cause any existing voids to expand uniformly

in all directions. In the fracture toughness test, the region ahead of

a sharp crack resembles the locally necked area of a tensile specimen.

Van Stone et al. 1 recently reviewed various models that have been

proposed to describe fracture toughness in terms of the dimpled rupture

fracture process. It is generally assumed that voids nucleate in the

plastic zone ahead of the crack tip at relatively low plastic strains and

that most of the energy measured in the J- integral test procedure is re- lated to void growth. 8

" 10 Van Stone et al. argued that any successful

modeling of ductile fracture ahead of a sharp crack must explicitly treat

all three stages of the process: void initiation, growth, and coalescence.

In this paper, we examine the fracture process in a series of seven

austenitic stainless steels broken at 4 K. A careful fractographic and

metallographic study was conducted of the broken specimens to observe the

dimpled rupture process. From these observations, a model for fracture

toughness was developed that relates the void nucleation process to the

mechanical properties that were presented in another report. 11

MATERIALS

The seven laboratory heats of steel were supplied in the form of

25-mm-thick plates in the hot-rolled condition. The plates were annealed

at 1170°C for 2 h and then water quenched. The treatment produced

coarse-grained, fully annealed, austenitic structures. The composition

grain size, A, and inclusion spacing of the seven alloys are presented ir. Table I. They differ mainly in their Ni and N contents.

115
Table I. Composition, Grain Size, and Inclusion Spacing of the Alloys

Alloy Content Grain

Number

Cr Ni Mn C

(wt.%)

N S P Si

size

(itm)

A

(itm)

1 18.8 5.6 4.1 0.020 0.256 0.005 0.014 0.25 122 47

2 18.4 8.9 4.2 0.018 0.281 0.005 0.014 0.28 165 63

3 20.5 12.8 5.5 0.024 2.265 0.004 0.015 0.34 200 66

4 20.8 14.9 5.2 0.028 0.277 0.004 0.015 0.41 160 63

5 18.7 8.7 3.7 0.014 0.093 0.005 0.013 0.23 220 57

6 19.7 11.5 3.9 0.012 0.141 0.005 0.013 0.27 250 57

7 20.8 14.7 4.4 0.016 0.197 0.004 0.014 0.33 207 70

The inclusion contents of these alloys were nearly identical.

Inclusion spacing was calculated by counting all of the inclusions larger

than 0.5 /xm in diameter in a given area, normalizing that number to

1 mm2

, and taking the inverse square root to obtain an approximate value

of spacing. Energy -dispersive x-ray analysis determined that the inclusions found in the polished cross sections of these alloys were predominantly MnSi and MnS types.

TEST PROCEDURES

Smooth, 6 -mm- diameter , 38 -mm- gage tensile specimens and 22 -mm- thick

compact specimens were machined from the annealed plates. Standard ASTM

E 8 and E 813 procedures were followed for obtaining the tensile and

fracture toughness properties of these alloys. Specific descriptions of

the techniques are available elsewhere. 7

Table II . Summary of Mechanical Properties at 4 K

Ni

content

N

content

a *ya *u

Of* e * In

A

o

A J Ic* Ki c (J)u Allov (wt.%) (wt.%) (MPa) (MPa) (MPa)

r Af (kJ/m2 ) (MPa * m )

1 5.6 0.26 1050 1418 1710 0.21 27 75

2 8.9 0.28 950 1650 3160 0.55 99 143

3 12.8 0.27 1000 1580 3130 0.78 201 204

4 14.9 0.28 1165 1630 2990 0.66 276 239

5 8.7 0.09 460 1417 2550 0.55 360 273

6 11.5 0.14 720 1370 2770 0.70 288 244

7 14.7 0.20 870 1425 2790 0.67 312 254

* Average value from 2 or 3 tests; <7y + au are engineering stress values,

of is the true stress at fracture, and (f is the true strain at fracture,

t Measured value from one test.

116
RESULTS

The results of tensile and fracture toughness tests at 4 K are listed in Table II. The variation of the Ni and N contents produced a wide range of mechanical properties. Nitrogen, dissolved into the fee

lattice at the interstitial sites, controls the yield strength of these

austenitic alloys at cryogenic temperatures. Yield strength affects the

fracture toughness, similar to the trend shown in Figure 1, but at a con- stant yield strength, Ni content determines the fracture toughness of

these alloys (see Figure 2) . Increasing Ni and N concentrations suppress

the strain- induced martensitic transformation, but the transformation

does not directly affect the Oy and Jjc mechanical properties. 11

Scanning electron microscopy was performed on all the test specimens

to determine the fracture morphology and the microstructural features

that participate in the fracture process. The observations, summarized

in Table III and illustrated in the next three figures , were always made

in the fracture -surface region with the highest triaxial stress. Most

specimens fractured in a ductile manner, which is characterized on a macroscopic scale by large shear lips on the tensile specimen and elastic

plastic behavior in the fracture toughness test. On a microscopic scale,

the predominant features were large dimples that originated at the MnSi

and MnS inclusions (see Figure 3) . These inclusions were spherical in

shape, between 1 and 2 /im in diameter, and identical to those found on

the polished cross section (see Table I)

.

NICKEL CONTENT, wt%

Figure 2. Fracture toughness vs. Ni

content for alloys 1 through 4 with constant N

content and similar yield

strengths

.

117
Table III. Summary of Fracture Behavior

Allov Tensile Fracture Toughness

1: 5 . 6Ni—0 . 26N 2-/im dimples;

small shear lips

facets

;

linear elastic

2: 8.9Ni—0.28N 20—30 /im dimples;

large shear lips

facets

;

elastic—plastic

3: 12.8Ni—0.27N 20—30 /im dimples;

large shear lips

30 -/im dimples;

elastic—plastic

4: 14. 9Ni—0 . 28N 20—30 /im dimples;

large shear lips

20—40 /im dimples;

elastic—plastic

5: 8 . 7Ni—0 . 09N duplex dimples,

2- and 10-/im;

large shear lips

duplex dimples,

2- and 10 -/im;

elastic—plastic

6: 11.5Ni—0.14N 20-/im dimples;

large shear lips

20—30 /im dimples;

elastic—plastic

7: 14. 7Ni-0 . 20N 20 /tm dimples;

large shear lips

20—30 /im dimples;

elastic—plastic

Specimens from alloys 1 and 5 exhibited a much smaller dimple size

(1 to 2 /im in diameter) on the fracture surface (see Figure 4). No inclusions are visible in the smaller dimples; however, the existence of

inclusions smaller that 0.5 /im in diameter, the smallest counted on the

polished cross sections, cannot be discounted. 12

Fracture toughness specimens of alloys 1 and 2 had a distinctly dif

ferent fracture morphology than the other specimens in this test series.

Their fracture surfaces were faceted; the facet size varied from 2-/im

steps on the surface to a relatively smooth surface extending across a complete austenite grain. This type of faceted fracture, shown in Fig- ure 5, has been described by Tobler and Meyn13 as a process of slip-band

cracking along {111} planes.

Selected tensile specimens were sectioned along the tensile axis

after testing to observe the micros true tural processes that led to the

formation of the large dimples (Figure 3) on the fracture surface. Fig- ure 6 shows a typical example. The rounded features of the dimples on

the fracture surface are visible at the edge of the cross section. The area directly below the fracture surface appears undamaged despite the

118
Figure 3. Frac to graph (SEM) of the tensile specimen for alloy 2 shoving

20- to 30-pm dimples in the center region of the fracture

surface

.

Figure 4. Fractograph (SEM) of the tensile specimen for alloy 1 show .

2-/im dimples on the fracture surface.

119
Figure 5. Fractograph (SEM) of the fracture toughness specimen for - _ alloy 1 showing facets on the fracture surface.

20/Ml

Figure 6. Micrograph (SEM) of the polished cross section of a tensile

specimen of alloy 3 showing void initiation at inclusions with

limited void growth.

120
high local stresses and strains present in the area. There is no gradient of voids below the fracture surface. Void area fraction vs. true

strain for one representative tensile specimen is shown in Figure 7. An

arbitrary function has been drawn through the data points to emphasize

the limited amount of void growth present below the tensile fracture

surface. A value of nucleation strain can be defined as the strain at zero area fraction of voids. In this example, the nucleation strain is

about 0.75 of the true strain to fracture of the specimen, but approximately equal to the average true strain to fracture reported in Table II

for the alloy.

Figure 7 . The area fraction of voids in the polished cross section of

a tensile specimen from alloy 3 vs. the applied true strain.

A cross section through the crack tip of a tested, but unbroken

compact tensile specimen from the highest yield strength alloy is shown

in Figure 8. The applied J at the crack tip was approximately equal to

J Ic (zero physical crack growth). The crack tip was not blunted significantly, and damage was accumulating in the plastic zone before actual

crack growth was detected. The damage was focused straight ahead, along

the path of subsequent macroscopic crack growth. This suggests that a stress-controlled process is important. 14

121
Load

Direction

Crack Growth _ Direction

End of | , 20 /im

Fatigue Crack J |

Load

Direction

Figure 8. Micrograph (SEM) shows a cross section through the crack

tip of a tested, but unbroken, compact tensile specimen

of alloy 4. The applied J is approximately equal to JjcDISCUSSION

Development of Model

The fracture process in these steels can be interpreted in terms

of two separate stages: 1) a nucleation stage and 2) a growth and coales- cence stage. In the uniaxial tensile test, the void nucleation stage occurs at strains very close to the fracture strain (Figures 6 and 7)

.

Growth and coalescence occur very rapidly and require nearly zero additional strain. The micrograph of the crack tip (Figure 8) shows a similar trend for limited void initiation at applied energy levels less than

J IcThese observations indicate that the conditions for nucleation are difficult to reach and require a large plastic strain, but once nucleation occurs, the conditions for growth and coalescence have already been

met so fracture can occur with little or no additional applied strain.

Broek15 describes a very similar type of behavior in high-strength aluminum alloys, in which the nucleation stage dominates the fracture process.

From the observations, we assume that the nucleation of microvoids

at the nonmetallic inclusions is the critical event that controls the

fracture of austenitic stainless steels at 4 K. LeRoy et al. 18 modeled

122
the nucieation condition in a tensile test specimen by assuming that

nucieation of a void at an inclusion occurs when the stress across the

interface exceeds a critical value (a c ) that is related to both the local

stress (a^, a function of the strain), and the macroscopic stress (r-) :

az " aL

+ am

and

where is the nucieation strain and itj} is a parameter that relates the

stress produced by dislocations piled up at an inclusion to the average

strain. We modified the equation of LeRoy et al . to use the full value

of the applied stress rather than the hydrostatic component of the applied stress for the macroscopic stress. Work of Cox and Low 1 " on highstrength steels shoved that void nucieation was dependent on the applied

stress rather than the triaxial stress.

In the fracture toughness test, the stresses and strains ahead of a stationary crack subjected to tensile (Mode I) opening can be predicted

by continuum mechanics. For elastic-plastic conditions, a solution for

the local strain (O at a distance ahead of a blunt crack has been calculated by Hutchinson, 18 and Rice and Rosengren 19 (HRR) in terms of the

J- integral (J)

3 6 <36 a I X

o o n

n/n+1

s

where a and n are constants of the Ramberg-Osgood fit to the material's

true-stress vs. true -strain curve, £ 0 and oQ are the yield strain and

stress in the Ramberg-Osgood fit, and r is the distance from the crack

tip. The term In , a function of n, is about 4.5, and « contains the

angular dependence of the strain.

The constant n is of the order of 10, so for simplicity, we assume n/(n+l) - 1. For crack advance all along the crack front, we assume that

the relevant distance, r, is the spacing of large inclusions, L. With

J - K

2 /E, where K is the stress - intensity factor and E is Young's modulus , then the HRR solution for strain ahead of a crack at the important

inclusions is:

K

2

6 “

4.5a EL

* o Nucieation was observed only very near the crack tip of the fract-rtoughness test specimens. Thus, nucieation strain is a major part c

: plastic strain needed to initiate tearing. Therefore, we equate the - . cleation strain to the strain that is one inclusion spacing ar.ead cf a blunt notch in the fracture toughness test and obtain an etuatirn : r stress - intensity factor associated with the nucieation process

123
(a - 3.5a ) ' c y

'•nuc 'H (4.5a EL)'

y

We have also assumed that the am - 3.5ay, because of crack- tip

constraint

.

( 6 )

The value of ac in this model is the cohesive energy of the interface between the defect and the matrix. We expect the value of ac to be

relatively independent of the chemical composition of the matrix, but it can depend on the size, shape, and type of inclusions. Exact values of

ac for these alloys are not known. LeRoy et al. 16 and Argon and Im20

obtained ac values from 1200 to 1800 MPa for cementite particles in a ferritic steel matrix. In general, it appears that the ac value for a tightly bond particle in a ductile matrix is about E/120. 20 In the terms

of our model for fracture toughness, this value must be multiplied by 3 because we used the applied stress rather than the hydrostatic component

of the stress in Equation 2. The elastic properties of austenitic stainless steels at cryogenic temperatures vary slightly, depending on the

exact chemical composition, but are essentially the same as ferritic

steels at room temperatures. 21 Therefore, we will assume that ac is E/40

or 5000 MPa for all the alloys.

The parameter ay in Equation 6 is also an unknown parameter that is a measure of the local stress exerted by the dislocation structures. An

idealized model would be of a slip band in the austenitic matrix where

dislocations have piled up against a barrier (a nonmetallic inclusion)

.

In this case, the ay constant represents the stress concentration ahead

of the pile-up and is a function of the stacking- fault energy. The dislocation structures in the vicinity of the inclusions in these steels are difficult to characterize completely, so ay cannot be calculated from

first principles. We assume that ay is a relatively strong function of

the Ni content and not the N content. We do expect dislocation mobility

and the ability to cross slip to vary with stress state, 22 so the value

of determined from the tensile test results cannot be used to predict

the fracture toughness

.

Evaluation of Model

To test this model for nucleation-controlled fracture toughness, we need to determine how ay varies with the Ni content. Therefore, considering only the highest strength alloys (1 through 4 in which Ni content

varies and N content is constant) , we set - Kjc

(J) and calculate

the change in ay from Equation 6 for the range of Ni values at a constant

yield strength using L (average value of 0.060 mm), ay (average value of

1040 MPa) , ac (estimated value of 5000 MPa) , and E (estimated value of

205 GPa) . The calculated values are plotted vs. Ni content in Figure 9.

Now we can evaluate the model and our assumptions. First, we con- sider alloys 5, 6, and 7 that have a range of both N and Ni contents. We can estimate ay from the Ni content in Figure 9 and calculate a Knuc us- ing the measured yield strengths and the constant values of L, ac , and E.

124
7

c/f 0

0

E

as I

5

4

3

2

C/5

C/5 05

CO

1

2 4 6 8 10 12 14 16

wt. % Nickel

Figure 9. The stress parameters ac and in the nucleation

controlled model vs. Ni content.

Table IV. Summary of Data Used in Nucleation Model

for Fracture Toughness

Alloy

aH

(MPa)

a

c (MPa)

ay

(MPa)

1 4350 5000 1050

2 2280 5000 950

3 1600 5000 1000

4 1365 5000 1165

5 2900 5000 460

6 1900 5000 722

7 1375 5000 865

K

nuc

(MPa-m1*)

*4 C

< J > (MPa • m’1

)

K /

73 75 0.98

169 143 1.18

221 204 1.08

173 239 0.72

186 273 0.6

260 244 1.07

318 254 1.25

125
We can take the same approach to examine the prediction of Knuc for

alloys 1 through 4, using the measured yield strengths rather than the

average. The variables used in the calculations and the results for all

seven alloys are found in Table IV.

/ The ratio of Knuc/Kjc

(J) for alloys 1, 3, and 6 is very close to 1, as expected, if nucleation does control the measured fracture toughness.

For alloy 5 with the lowest yield strength, the model underpredicts the

measured value by 30%. We interpret this to be a yield strength effect.

For lower strength alloys, we expect nucleation strains to be less than

the fracture strains so that additional plastic strain is required to

grow the voids to the point of coalescence. For alloys 2, 4, and 7, the

ratio is not very close to 1, indicating that the model is sensitive to

the values of the unknown parameters, ajj and ac

. If there is a variation

in the ac value with Ni or N or Ni-plus-N content or if the true variation of ay is different from that shown in Figure 9, then the model

cannot accurately predict the fracture toughness.

Figure 10. Fracture toughness, model prediction and measured,

for Fe—18Cr—lONi alloys vs. yield strength. 8

126
A second, independent test of the model is found in Equation 1 for

the Fe-18Cr-10Ni alloys with various N contents. From Figure 9, we es- timate <7^ to be 2400 MPa and assume the same values for ac and E. The

average inclusion spacing for this alloy series was 0.055 mm. The predicted values of Knuc and the calculated KIc (J) from Equation 1 are plotted as a function of the yield strength in Figure 10.

The model predicts a maximum toughness at a yield strength of 400 MPa

and a decrease in toughness for higher yield strengths. The slope of the

predicted curve for yield strengths greater than 600 MPa matches that for

the experimental data. The different intercept for the predicted values

from that of the model and measured values from Equation 1 could be due

to a systematic error. In general, the model works reasonably well for

the higher strength alloys in which we expected nucleation to dominate

the fracture process. For yield strengths below 600 MPa, the model underpredicts the fracture toughness, reflecting the lesser role of nucleation

and a significant contribution of void growth and coalescence to the mea- sured fracture toughness

.

SUMMARY

In high-strength austenitic stainless steels fractured at 4 K, large

plastic strains, nearly equal to the true fracture strain, are needed to

nucleate voids in the tensile test. In the fracture toughness test, the

void nucleation is limited to the region directly ahead of the advancing

crack tip. On the basis on these observations, we assumed that tearing

in the fracture toughness test is a stress -controlled process and begins

when the stress at an inclusion-matrix interface ahead of the crack

reaches the critical value required for nucleation. From this criterion

for ductile fracture, we have developed a model that relates the stressintensity factor for nucleation to the inclusion spacing, the critical

stress for void nucleation, the material's strength, and the local stress

at inclusions produced by dislocation pile-ups. The model can predict

reasonably well the relative effect of changes in the dislocation structure and yield strength on the fracture toughness.

The model represents an attempt to understand fracture toughness on

the basis of the micromechanisms that control void nucleation. We have

assumed that the fracture toughness represents an average nucleation

strain. The model can be used to explain the results for high-strength

austenitic steels at low temperatures where the inclusions are widely

spaced in, and tightly bound to, the matrix. The model has shortcomings:

it considers only the nucleation contribution to the fracture process; ic

is sensitive to the stress terms, so they must be known; and it ignores

the effect of inclusion size, shape, and local spacing. The model has

many advantages: it relates the effect of dislocations and interfacial

stresses directly to the fracture toughness, and it does not require a detailed analysis of the fracture surface features, which we think are more representative of the tearing modulus than the fracture toughness of

127
high-strength materials. The development of a more comprehensive model

is certainly possible through more extensive testing and thorough material evaluation.

ACKNOWLEDGMENT

This work was partially supported by the U.S. Department of Energy,

Office of Fusion Energy. The seven laboratory heats of steel were supplied by the E. 0. Paton Institute of Electrowelding, Kiev, USSR.

REFERENCES

1. R.H. Van Stone, T.B. Cox, J.R. Low, Jr. , and J.A. Psioda: International Metals Review, vol. 30, 1985, p. 157.

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1980, p. 219.

3. J.M. Krafft: Applied Materials Research, vol. 3, 1964, p. 88.

4. G.T. Hahn and A.R. Rosenfield: Metallurgical Transactions, vol. 6A,

1975, p. 653.

5. J.P. Hirth and F.H. Frees: Metallurgical Transactions, vol. 8A,

-1977-r p. 1165.

6. R.J.H. Wanhill: Engineering Fracture Mechanics, vol. 10, 1978,

p. 337.

7. 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.

8. J.R. Rice and M.A. Johnson: in Inelastic Behavior of Solids,

M.F. Kanninen, W.G. Adler, A.R. Rosenfield, and R.I. Jaffee, eds.,

McGraw-Hill, New York, 1970, p. 641.

9. R.O. Ritchie and A.W. Thompson: Metallurgical Transactions, vol 16A,

1985, p. 223.

10. W.M. Garrison, Jr.: Metallurgical Transactions, vol, 17A, 1986,

p. 669.

11. R.P. Reed, P.T. Purtscher, and K.A. Yushchenko: in Advances in Cryogenic Engineering—Materials , vol. 30, R.P. Reed and A.F. Clark,

eds., Plenum Press, New York, 1986, p. 43.

12. A.W. Thompson and P.F. Weihrauch: Scripta Metallurgica, vol. 10,

1976, p. 205.

128
13. R.L. Tobler and D. Meyn: in Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures—XI , R.P. Reed, ed.

,

National Bureau of Standards, Boulder, Colorado, 1988, p. 89.

14. R.M. McMeeking: Journal of Mechanics and Physics of Solids, vol. 25,

1977, p. 357.

15. D. Broek: Engineering Fracture Mechanics, vol. 5, 1973, p. 55.

16. G. LeRoy, J.D. Embry, G. Edwards, and M.F. Ashby: Acta Metallurgica

,

vol. 29, 1981, p. 1509.

17. T.B. Cox and J.R. Low, Jr.: Metallurgical Transactions, vol. 5A,

1974, p. 145.

18. J.W. Hutchinson: Journal of Mechanics and Physics of Solids,

vol. 16, 1968, p. 13.

19. J.R. Rice and G.K. Rosengren: Journal of Mechanics and Physics of

Solids, vol. 16, 1968, p. 1.

20. A.S. Argon and J. Im: Metallurgical Transactions, vol. 6A, 1975,

p. 839.

21. H.M. Ledbetter: in Austenitic Steels at Low Temperature, R.P. Reed

and T. Horiuchi, eds., Plenum Press, New York, 1983, p. 83.

22. S. Floreen, H.W. Hayden, and T.M. Devine: Metallurgical Transactions-

, vol. 2A, 1971, p. 1403.

129

LOW- TEMPERATURE PHASE AND MAGNETIC INTERACTIONS

IN FACE -CENTERED -CUBIC Fe-Cr-Ni ALLOYS*

C. Almasan, T. Datta, R. D. Edge, and E. R. Jones

University of South Carolina

Columbia, South Carolina

J. W. Cable

Oak Ridge National Laboratory

Oak Ridge, Tennessee

H. Ledbetter

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

_The magnetic properties of a set of nine isostructural

face -centered- cubic Fe—Cr—Ni alloys were studied by SQUID magnetometry, neutron diffraction, and ultrasonic techniques from

5 to 300 K. Type-1 antiferromagnetic ordering was observed

below the Neel temperature (T^) . The dc susceptibility [x(T)

]

did not exhibit a simple Curie—Weiss dependence. Above Tjj,

x(T) contained a temperature -independent component (x 0 ): *(T) - Xo + C/(T + 9). The lattice parameter (a) systematically

influenced Tjj, which decreased from 47.9 ± 0.05 to 35.0 ± 0.5 K

as a increased only 0.25%. The average magnetic moment, “ 0.5 /xg, obtained from neutron scattering was lower than the * 1.0 /ig value obtained from the SQUID measurements. Meanfield estimates of the antiferromagnetic first-neighbor ex- change interaction (J x ) and the ferromagnetic second- neighbor

interaction (J 2 ) indicated that |J 2 /Ji| “ 1.5. We take this as evidence for RKKY interaction and argue self -consistently that

only the external d electrons are responsible for the localized

average moment. This may mean that s-d hybridization of the

external electrons is weak in these alloys.

*Draft

;

intended for Physical Review.

131
INTRODUCTION

The nature of the ground state of 7 -Fe, Fe in the face-centeredcubic (f.c.c.) structure, has been controversial ever since Weiss and

Foex argued that 7 -Fe has an antiferromagnetic ground state , 1 ’ 2 unlike

the familiar body-centered-cubic a-Fe. Two other issues 3 ’ 4 related to

itinerant-electron magnetism are also important in 7 -Fe: (i) the nature

of the magnetic interaction and (ii) the localization of the magnetic

moments

.

Although the f.c.c. allotrope is unstable at room temperature, it

may be obtained by alloying with Ni and then stabilizing the resulting

structure by adding Cr, Mn, V, or Cu. The behavior of Fe—Cr—Ni alloys is

typical of iron in the f.c.c. structure . 1 These alloys are the basis for

commercial "nonmagnetic" austenitic stainless steels. Many of their use- ful properties relate intimately to their magnetic behavior6 ’ 6 because

the magnetic properties reflect the electronic structure of the alloys.

This system is also a solid-state experimental material with a rich array

of physical properties (including composition and lattice spacing) that can be systematically varied. Clustering and the amount of short-range

order may also be influenced by careful thermal—mechanical treatment.

This study concerns the low- temperature (from 5 to 300 K) properties

of a set of austenitic Fe—Cr—Ni alloys with Ni concentration less than

10 at.%. Emphasis is given to ultrasonic measurements, the effects of

lattice dilation on the phase transition, the exchange interaction, and

the average magnetic moment as determined by SQUID magnetome try , and

neutron scattering.

EXPERIMENTAL DETAILS

Materials

All nine sets of specimens studied were type-304 polycrystalline

austenitic stainless steel. As mentioned above, these isostructural

alloys represent a close realization of 7 -Fe. The lattice spacing (a) was systematically dilated by introducing of interstitial C and N atoms.

In alloys 1 through 9, the C+N concentration ranged from 0.5 to 1.5 at.%.

The crystal structures and values of a were determined at room temperature for all nine compositions by using a modified commercial x-ray

spectrometer. For four samples, the values of a at different temperatures were determined from the nuclear Bragg peaks observed in the

elastic-neutron-diffraction powder pattern. Both microprobe (EDAX) and

chemical analysis were performed to determine the chemical composition

of the specimens.

Ultrasonic Measurements

Acoustic velocities were measured as functions of polarization and

temperature by a technique detailed elsewhere . 7 ’ 8 Briefly, a pulse-echo

method was used to obtain the velocity, v - 2 SL/r

,

where r denotes the

132
round-trip time for an ultrasonic pulse between two flat, parallel faces

of the specimen separated by a distance 2 . Both longitudinal and transverse waves were studied.

Magnetic Measurements

Specimens were cut into rectangular prisms with a typical size of

1.5 mmx 3 mmx 7 mm. They were cleaned in a solution containing equal

parts of HCl and HN0 3 and a small amount of H2 S0 4

. Ultrasonic agitation

accelerated the cleaning. Finally, they were washed in distilled water,

rinsed in alcohol, and dried in air. The specimens were initially cooled

to 5 K in zero applied magnetic field (B)

.

With a computer -controlled SQUID variable- temperature susceptometer, 9 the longitudinal dc susceptibility [x(T)j was measured from 5 to

200 K in magnetic fields of 1, 5, 20, and 50 mT. Thermohysteresis was not observed, although it has been reported in related alloys. 10 ' 12

Neutron-Diffraction Measurements

Specimens for the elastic-neutron-scattering experiments were flat,

2-cm x 2-cm x 0.5-cm polycrystalline plates. A closed-cycle system

controlled the specimen temperature; measurements were made from 10 to

about 300 K. A thermal neutron beam from the High Flux Isotope Reactor

with a wavelength of 0.2350 nm (14.816 meV) was used.

Experiments were performed on the HB1A triple -axis neutron spectrometer adjusted for elastic scattering. A pyrolytic graphite filter

removed the A/2 components. We obtained low-resolution measurements of

angular powder patterns (5° < 29 < 80°) and high- resolution angular scans over the anticipated regions of the magnetic and nuclear Bragg peaks as

a function of temperature.

At all temperatures studied, large nuclear peaks arose from the

(111) and (200) lattice reflections. The widths of these nuclear lines

were limited by the spectrometer resolution, which indicated a welldefined long-range crystalline order. A (110) peak was also observed.

The neutron counts for the (110) peak were composed of a temperature - independent nuclear component and a temperature -dependent component

caused by magnetic scattering. The small critical-scattering component

was observed only above the Tjj appropriate for each specimen. Below T»;, the magnetic scattering increased with decreasing temperature. Because

A/2 components were eliminated from the measurements, we think that onlv

a small fraction of the (110) line arises from A/2 contamination. The

small (001) peak reported by Ishikawa et al. 1 for a monocrystal specimen

of similar composition was not observed with the polycrystalline sample,

even after carefully scanning the anticipated angular region.

A large, approximately uniform background arose from incoherent or diffuse scattering in the specimen. The ratio of the magnetic peak

height to this background was typically 0.2, although large differences

occurred among the specimens

.

133
Table I. Chemical Composition (at.%)

Alloy Fe Cr Ni Mn P S Si Ho Cu C N C+N

1 68.19 19.4 9.49 1.58 0.037 0.033 0.98 0.115 0.182 0.174 0.165 0.340

2 68.19 19.4 9.49 1.58 0.037 0.033 0.98 0.115 0.182 0.285 0.169 0.454

3 68.19 19.4 9.49 1.58 0.037 0.033 0.98 0.115 0.182 0.514 0.173 0.687

4 68.65 19.1 9.30 1.50 0.039 0.029 1.12 0.118 0.171 0.252 0.539 0.791

5 68.65 19.1 9.30 1.50 0.039 0.029 1.12 0.118 0.171 0.349 0.531 0.88

6 68.65 19.1 9.30 1.50 0.039 0.029 1.12 0.118 0.171 0.440 0.507 0.94

7 66.32 21.5 9.34 1.42 0.043 0.027 1.08 0.112 0.173 0.165 1.13 1.29

8 66.32 21.5 9.34 1.42 0.043 0.027 1.08 0.112 0.173 0.284 1.09 1.38

9 66.32 21.5 9.34 1.42 0.043 0.027 1.08 0.112 0.173 0.435 1.04 1.4

Fig. 1. Lattice parameters determined by x-ray and

neutron diffraction at room temperature.

134
RESULTS

Structure and Composition

The chemical compositions determined for the specimens appear in

Table I. The interstitial concentration (C^) clearly correlates with the

lattice spacing (a) at room temperature (Fig. 1) . We determined the rate

of increase in lattice spacing per atomic percent of interstitial C+N to

be da/dCi - 8.62 x 10' 3 A C^1

.

Elastic Constants

The ultrasonic experiments 7 ’ 8 indicated that the elastic constants

were strongly affected by cooling the specimens through the paramagnetic—

antiferromagnetic transition temperature. The bulk modulus softened

above and around T^j, but the shear modulus softened only at and below T^j.

Also, in the low- temperature region, an anomalous temperature dependence

was observed in Poisson's ratio. These observations are consistent with

a paramagnetic—antiferromagnetic phase transition at Tjj.

Magnetic Susceptibility

The magnetic susceptibility [x(T)] of each of the nine alloys showed

a distinct peak as a function of temperature. This behavior is characteristic of a paramagnetic—antiferromagnetic phase transition. However,

X(T) was independent of the Curie—Weiss temperature. Instead, over a limited range of temperatures above the transition, the susceptibility

was well described by

x(T) - Xo + C/(T - 0). (1)

Here, Xo * ar*d 9 are independent of temperature. The constants C and

9 represent the Curie constant and the paramagnetic Curie temperature,

respectively. The temperature independence of Xo may result from the

underlying diamagnetism, Pauli spin paramagnetism, 13 and a saturated

ferromagnetic component caused by small ferromagnetic clusters within the

alloy. 14 Table II gives the values of Xo » C, Tjj, and /i e ff (the effective

magnetic moment per atom) obtained from the analysis of Eq . (1). The

small and negative values of 9 for all specimens (Table III) indicate

a weak net antiferromagnetic interaction. The value of Xo generally

decreases as the lattice parameter increases. This behavior is consistent with expected changes in the Landau—Peier Is and Pauli susceptibilities 13 ’ 14 caused by the lattice dilation and the reduction of the Fermi

energy.

135
Table II. Summary of magnetic measurements

Alloy

* Xo *

C 4 effective (d+s) Me ff* Meff^

(10' 7 m

3 /kg) (10" s K-m3 /kg) (K) electrons/atom (mb) (mb)

1 3.68 ± 0.15 3.49 ± 0.19 47.9 7.6 1.11 0.63

2 2.89 ± 0.05 3.64 ± 0.09 48.1 7.6 1.13

3 2.70 ± 0.06 3.5 ± 0.09 47.1 7.6 1.11 0.75

4 2.38 ± 0.10 3.19 ± 0.15 44.3 7.67 1.06

5 2.72 ± 0.06 3.66 ± 0.10 43.8 7.67 1.13

6 2.13 ± 0.21 4.47 ± 0.31 42.25 7.67 1.25

7 2.92 ± 0.03 4.17 ± 0.07 36.5 7.85 1.21 0.48

8 2.17 ± 0.01 3.05 ± 0.05 36.4 7.85 1.04

9 1.45 ± 0.61 4.13 ± 0.68 35.0 7.85 1.21 0.66

Constants obtained from a fit of x - Xo + C/(T + 0) to the data

tordering temperature obtained from susceptibility maxima

^ Average magnetic moment per atom, ^e ff, as resultsdetermined by static magnetic

'Ueff

- determined from neutron- diffact ion data

Table III, Calculated parameters

Alloy 0

(K) 7i/7 2 -0/Tn

J 2 /Ji 7

1

(kg/m3

) Ji/k

(K)

1 -13.96 -0.699 0.291 -1.398 -13.289 -0.824

2 -13.15 -0.714 0.273 -1.428 -12.632 -0.783

3 -14.86 -0.680 0.315 -1.360 -13.268 -0.822

4 -14.77 -0.667 0.333 -1.334 -15.223 -0 . 944

5 -13.61 -0.684 0.311 -1.368 -11.768 -0.729

6 -6.10 -0.832 0.144 -1.664 -8.123 -0.504

7 -4.20 -0.862 0.115 -1.724 -3.186 -0.454

8 -7.01 -0.785 0.193 -1.570 -10.69 -0.663

9 -5.75 -0.812 0.164 -1.624 -7.406 -0.459

136
The maximum value of the susceptibility decreased with increased

dc magnetic field. This behavior is similar to that reported 15 for a 20-Cr, 16-Ni (wt.%) alloy. In the present study, however, the peak was noticeably sharper and, at the fields used, did not disappear as in the cases reported by Warnes and King 15 and by Ishikawa et al. 1 Over the

fields studied (1 to 50 mT) , the position of the susceptibility maximum,

which we take as Tjy, was unaffected by the field magnitude. In contrast,

Warnes and King reported a gradual reduction in Tjj with increasing magnetic field for fields from 58 to 460 mT.

Pinski et al. 4 predicted that increasing the lattice spacing decreases the transition temperature. We observed that Tjj decreased from

47.9 ± 0.5 K for the first alloy to 35.0 ± 0.5 K for the ninth alloy.

Since the monotonically varying quantity in these alloys is the lattice

parameter (which increases from alloy 1 to alloy 9) , we think that the

observed change in Tjj is caused by interdependence of the lattice parameter and the transition temperature. This shift in Tjg for three alloys is

shown in Fig. 2.

Fig. 2. Magnetic evidence for the systematic shift in the transition

temperature. The specimens were first cooled to 5 K under zero applied magnetic field. These measurements were made during

warming in a 1.0-mT field.

Neutron Diffraction

A typical neutron-diffraction powder pattern at 10 K is shown in

Fig. 3. For quantitative analysis of the neutron-diffraction peaks, the

observed background was eliminated by numerical methods. For the broad

peaks, an average of several high- temperature (=* 300 K) spectra, taken over the same angular region, was subtracted from the low- temperature

spectrum being analyzed. This procedure provided a sharp, single peak

137
that was then fitted to a Gaussian profile. For measurements showing an initial sharp peak, the background was assumed to be a second- order polynomial, and the spectra were fitted with a Gaussian profile and the background. Results of these methods generally agreed. The results of the

analysis of the magnetic lines for two specimens are shown in Fig. 4.

The temperature dependences of the full width at half maximum of the

lines and of the estimated coherence lengths are given in the inset.

The transition temperatures (Tjj) were also obtained from the

neutron- diffraction measurements by analyzing the magnetic -scattering

intensity (N) for temperatures less than Tjj. When N was scaled as N oc (Tjg — T) 2 ^, 16 we obtained 0.29 < (3 < 0.43. In all cases, the correlation coefficient exceeded 0.97, which indicated a good fit to the scaling law. Absolute structure factors for the (110) magnetic reflections

were obtained by means of an internal calibration against the nucleus

reflections. From these, the effective ordered magnetic moments were

determined.

Ishikawa et al. 1 showed that the presence of the (110) reflection

and the absence of (001) reflection in these alloys are evidence for co- linear magnetic structure, or type-1 antiferromagnetism. In this structure, successive planes of atoms perpendicular to the (001) direction

have atoms parallel and antiparallel to this axis. Such an arrangement

would not produce a (001) reflection because only those moment components

perpendicular to the scattering vector are effective in the scattering.

138
We observed the (110) peaks but found no measurable (C01) reflection for

any specimen studied; we conclude that the moments are either parallel

or antiparallel to the c axis. Hence, in all specimens studied, the

neutron- di-f fraction powder patterns are consistent with a low- temperature

type-1 antiferromagnetism.

Fig. 4. Temperature dependence of a magnetic peak for two alloys

The x

' s denote specimen 1; filled circles denote speci.T.en • Inset shows the coherence length and the full-width at halfmaximum as functions of temperature for specimen 1. 71. e Neel

temperatures, determined by magetometry, are also indicate

139
ANALYSIS

The values of Tjj obtained from the longitudinal ultrasonic wave

velocities and from the neutron-diffraction results were higher than the

values obtained from the susceptibility maxima; the values of Tjj obtained

from the transverse wave velocities were slightly lower than the values

from the susceptibility maxima. The general trend of Tjj versus lattice

parameter from both the elastic and the neutron results parallels that

of the magnetic -susceptibility measurements. A fit of the ultrasonic

results to a linear regression gives slightly different slopes for the

transverse and longitudinal wave velocities, but within their statistical

uncertainties, they agree with the slope obtained from the susceptibilities. A coherent picture of the low- temperature phase transitions in

these nine alloys is shown in Fig. 5. It is a composite of the effect of

the lattice expansion on Tjj as determined by SQUID magnetometry , longitudinal and transverse elastic anomalies, and the neutron measurements.

Overall, the results concur. The decrease in the transition temperature

as the lattice expands is clear.

UJ

QC3

I-<

cr

in

CL2

UJO

<0

Z

<

crH

55

50 -

45 -

40

35

30

1 —l

> O

\

\ °! G

\

a / O

i

\

o

o </ ^ a ^ \ c

Va \o\ O O * ^

A \ o

* \

\

\

T) Longitudinal sound velocity a Transverse sound velocity \

o Magnetic susceptibility \

O Neutron diffraction * A \

i l

> 3.588 3.592 3.594

UNIT-CELL SIZE (A)

3.600

Fig. 5. Composite picture of T^ versus lattice parameter. Measurements

of Tfj from static magnetometry, longitudinal and transverse

sound velocities, and neutron diffraction are shown.

140
II)

III)

'II
Calculations appropriate for the f.c.c. structure with type-1 antiferromagnetic ordering were made (Table III). 17 For all specimens, the value

of 0/Tjj lies between —0.33 and —0.12. The small values of 0 /Tn for these

antiferromagnetic alloys are in stark contrast with the large values (>1)

observed in the spin- glass alloys with similar compositions. 10 ’ 12 The

ratios of 9/Tc > 1 indicate strong "frustration" so that, even for a large net interaction 0, the spin-glass ordering temperature is suppressed by the frustrations. Conversely, in the alloys of this study,

the frustration effects were minimal.

Smart 17 showed that type-1 antiferromagnetic phase is possible over

this region of the ( 71 /72 - ©/Tjj) parameter space. This is a theoretical

consistency check of the presumed magnetic ground state and the experimental measurements

.

The negative values of J : and J 2 /J x in all cases indicate that the

first-neighbor interaction is antiferromagnetic, and the second-neighbor

interactions are ferromagnetic. Furthermore, |J 2 />li| > 1 implies that

the second-neighbor interaction is stronger than the first-neighbor

interaction. The larger value of J 2 compared with that of Jj does not

affect the effective-field approximation used here because there are twelve first neighbors, each contributing -|J X | towards 0, whereas there

are only six second neighbors, each contributing about 1 . 5J x

. Hence, 0

is still small and negative, or antiferromagnetic, consistent with ex- perimental observations.

Because of the deviations from the Curie—Weiss susceptibility and

because |J 2 /J X | > 1 , we conclude that the alloys of this study must have

strong itinerant-electron antiferromagnetic contributions. Perhaps some

of the external electrons are localized and the rest are in extended

states. The delocalized electrons contribute to the temperatureindependent susceptibility (xo) an<^ mediate the interaction of J and

the localized moments. Friedel oscillation would cause J to oscillate

as a function of the distance (r) . For large x, where x - 2r-kp and

kp is the Fermi wave vector, J is represented by the well-known RKKY

function: 19 > 20

J(x) xcosx — sinx

( 2 )

In this equation, J 0 is assumed to be independent of kp and r. Equation (2) may be used to determine the interactions J x and J 2 at r x - 0.707a and r 2 - a, respectively, for given values of kp. Instead,

we used Eq . (2) to determine the free-electron band-filling factor

[<t> - kp(?r/a)] for the nine alloys from the values of a and the previously determined J x and J 2

. At least in the mean-field limit, all the

experimental facts (type-1 antiferromagnetism, 8, Tjj, C, and a) were

self -consistently incorporated into the determination of the Fermi wave

vector. Equation (2) was solved numerically to determine the value of

142
kp, and hence, 0 , needed to satisfy J - -J x at r - r x and J - J 2 at

r - r 2 . The values of 0 for different alloys were: <£(1) ” 1.905,

0(2) - 1.905, 0(3) - 1.90, 0(4) - 1.90, 0(5) - 1.0904, 0(6) - 1.925,

0(7) - 1.926, 0(8) - 1.917, and 0(9) - 1.92. On the other hand,

kp - (3ir 2 n) x / 3

, (3)

where n is the density of free electrons. Considering that there are four atoms per unit cell and that ne ff is the effective free -electron

contribution per atom, Eq. (3) takes the form

keff “ (3*) l/3

1/3

(4)

When the expression for the wave vector as a function of the freeelectron band-filling factor, kp - 0(ff/a), is combined with Eq. (4), we obtain the following expression for ne ff.

ne ff * (*/12)0 3

. (5)

The value of ne ff for each of the nine alloys was calculated from

Eq. (5) and the band-filling-factor results. For all alloys, we obtained

ne ff * 2 electrons per atom. The total number of external d and s electrons calculated directly from the chemical composition gives ne ff =* 8.

When only the d electrons on the Fe and Ni atoms are counted, we obtain

ne ff =5. On the other hand, when only the s electrons are included, we obtain ne ff =* 2. This suggests that the observed exchange interaction,

the transition temperature, and the type-1 antiferromagnetic ordering can

be accounted for by considering the extended free-electron cloud caused

by only s electrons.

The observed sensitivity of Tjj to the lattice parameter (a) may be

qualitatively explained by the variation of the RKKY interaction with x. If x, the Fermi wave vector times the intemeighbor separation, is such

that there is a node between x x - (kp-a/72 ) and x2 - (kp-a), then the

first-neighbor and second-neighbor interactions (J x and J 2 ) will be oppo-

&2site. Furthermore, as shown in Fig. 6, if the values of x are close to

a node, then a change in a will greatly affect J x and J 2 in opposite

directions. So, a small increase in x x and x 2 will decrease |jj - J ; | by

a large amount and lower the transition temperature.

We cannot justify using Eq. (2) for the concentrated alloys studied

here because Eq. (2) is thought to be appropriate for interactions in a dilute system, but it has been reported21 that the behaviors of concen- trated and dilute alloys are often similar. So, empirically, it may be

appropriate to express the interaction by Eq . (2).

143
Fig. 6. A pictorial explanation of the extreme sensitivity of the

net interaction (Ji~

J

2 ) to the variation of x - kpr.

CONCLUSIONS

We studied a family of isostructural Fe—Cr—Ni alloys with the y-Fe

structure. Elastic, magnetic, and neutron- scattering measurements were

analyzed using self-consistent methods to determine the nature of the

magnetic ground state and the mechanism of the magnetic interaction. We

think that both localized and delocalized electronic contributions are present.

• Our estimates of kp suggest that only the s states are strongly

delocalized. At this stage, we have no explanation for the weak s—

d

hybridization, and we do not know how many of the d states are delocalized.

• It appears that the type-1 antiferromagnetic ground state is attained

by the ordering of the moments of the partially localized d orbitals.

These moments were modeled to interact via the extended s states.

This oscillatory RKKY interaction gives rise to a weak net antiferromagnetic behavior.

• The sharp decrease in Tjj with increasing electronic density or lattice

separation can be explained by the rapid change in the exchange interaction near the zeros of the RKKY function.

144
• For 7 -Fe alloys with fewer than 7 electrons per atom, we argue that

an antiferromagnetic ground state is to be expected, and for more than

8 electrons per atom, a ferromagnetic ground state is favored. 1 For

intermediate electron densities, unusual ordering, such as spin- glass

behavior, may be favored. Studies on monocrystals of these alloys

with this intermediate composition are in progress and will be re- ported later.

ACKNOWLEDGMENTS

The work at South Carolina was partially supported by USC grant

13070-K101 and by O.R.A.U.S. contracts to travel to Oak Ridge National

Laboratory for the neutron measurements. Work at Oak Ridge was supported

under contract DE-AC05 - 840R21400 with the U.S. DOE. At the National

Bureau of Standards, studies were supported partly by the DOE Office of

Fusion Energy. C. A. acknowledges a fellowship from Zonta International.

REFERENCES

1. Y. Ishikawa, Sci. Rep. Tohoku Univ . 58, 151 (1975); Y. Ishikawa,

Y. Endoh, and T. Takimoto, J. Phys . Chem. Solids 31, 1225 (1970);

Y. Ishikawa, M. Kohgi, and Y. Noda, J. Phys. Soc . Jap. 39, 675

(1975); Y. Ishikawa, in Physics and Application of Invar Alloys, ed.

H. Saito, Honda Memorial Series on MH. Sc. #3, Tokyo (1978).

2. A. Z. Menshikov, P. Burlet, A. Clamberod, and J. L. Tholence, Solid

State Commun. 39, 1093 (1981).

3. F. Gautier, in Magnetism of Metals and Alloys, ed. M. Cyrot (NorthHolland, New York, 1982), p. 1.

4. F. J. Pinski, J. Stanton, B. L. Gyroffy, D. D. Johnson, and G. M.

Stocks, Phys. Rev. Lett. 56, 2096 (1986).

5. A. K. Majumdar and P. V. Blanckenhagen, Phys. Rev. B 29, 4079

(1984).

6. R. M. Bozorth, Ferromagnetism (Van Nostrand, Princeton, New Jersev

1951), p. 148.

7. H. M. Ledbetter. Metall. Trans. A 11, 543 (1980).

8. E. W. Collings and H. M. Ledbetter, in Metal Science of Stainless

Steel, eds . E. W. Collings and H. W. King (Metallurgical Society,

New York, 1979) , p. 22.

9. #VTS-805, BTI(SHE) Corp. San Diego, California.

145
10. T. Datta, S. D. Levine, D. Thornberry, and E. R. Jones, Phys . Status

Solidi (b) 121, K125 (1984).

11. T. Datta, D. Thornberry, E. R. Jones, and H. M. Ledbetter, Solid

State Commun

.

52, 515 (1984).

12. T. Datta, D. Thornberry, C. Almasan, and E. R. Jones, Solid State

Commun. 56, 523 (1985).

13. R. M. White, Quantum Theory of Magnetism (Springer, Berlin, 1983).

14. C. Kittel, Introduction to Solid State Physics (Wiley, New York,

1976) .

15. L. A. A. Warnes and H. W. King, Cryogenics 16, 473 (1976); ibid 16,

569 (1976); ibid 21, 729 (1981).

16. A. Tucciarone, H. Y. Lau, L. M. Corliss, A. Delapalme, and J. M.

Hastings, Phys. Rev. B 4, 3206 (1971).

17. J. S. Smart, Effective Field Theories of Magnetism (W. B. Saunders,

Philadelphia, 1966).

18. E. R. Jones, T. Datta, C. Almasan, D, Edwards, and H. M„ Ledbetter,

Mater. Sci. Eng. 91, 181 (1987).

19. M. A,- Ruderman and C. Kittel, Phys. Rev. 96, 99 (1954).

20. T. Kasuya, Prog. Theor . Phys. 16, 45 (1956); K. Yosida, Phys. Rev.

106, 893 (1957).

21. G. Aeppli, J. J. Hauser, G. Shirane, and Y. J. Uemuria, Phys. Rev.

54, 843 (1985).

146
EFFECTS OF GRAIN SIZE AND COLD ROLLING

ON CRYOGENIC PROPERTIES OF COPPER*

R. P. Reed, R. P. Walsh, and F. R. Fickett

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The effects of grain size and cold rolling on the tensile

properties and electrical resistivity at 295, 76, and 4 K were

studied for oxygen-free, high- conductivity copper. Tensile

yield and ultimate strengths increase linearly with increasing

d

_i

* (d - grain diameter) , following the Hall—Petch relationship. At low temperatures, the dependence on grain size in- creases. Increasing grain size lowers resistivity slightly at

all temperatures. Cold rolling to 10 percent reduction of area

significantly increases the yield strength at all temperatures;

subsequent rolling produces smaller strength increases. Resistivity increases with cold rolling.

INTRODUCTION

In the annealed condition, copper is usually not strong enough for use as the stabilizing element for low- temperature superconductors or as a high- field magnet conductor. It must be strengthened or reinforced.

Two conventional ways to strengthen copper are cold working and reduction

of grain size. Yet quantitative low- temperature property data on the

effects of these metallurgical variables are not available. This paper

presents the effects of grain-size variation and cold rolling on the tensile properties and electrical resistivity of copper at 295, 76, and 4 K. The material studied was the oxygen- free, high -conductivity , silverbearing copper (99.95 Cu + Ag) designated by the Copper Development Association (CDA) as CDA 104.

In Advances in Cryogenic Engineering — Materials

,

vol. 34, Plenum,

New York, 1988, pp. 299-308.

147
EXPERIMENTAL PROCEDURES

Tensile Measurements

Tensile specimens taken from the plate material were rods; their

reduced sections were 6.3 mm in diameter and 38 mm in gage length, with

12 . 7-mm-diameter threaded ends. For rolled plate materials, the specimens were oriented so that the tensile force was applied transverse to

the rolling direction.

For all tests, specimen strain was monitored with three clip-on

strain-gage extensometers , 25 mm in gage length. They were mounted on

the specimen within the 38-mm gage length. From specimen to specimen,

the outputs of these extensometers reproduced within 14 nm with a sensitivity of 0.03 strain per volt (corresponding to a readout of ±0.04 /im)

.

The tensile testing machine had a cryostat fixture mounted to the

bottom of its screw-driven crosshead. The specimen strain rate was 8.8 x 10 s" 1

. Cryostat equipment and procedures using liquid helium and

nitrogen have been described. 1

Resistivity Measurements

Resistance measurements at room temperature (295 K) , in liquid

nitrogen (76 K) , and in liquid helium (4 K) were made on tensile specimens.- A conventional four-probe resistivity technique 2 with a commercial

nanovoltmeter and an automated data acquisition system was used. Voltage

taps were placed along the uniform section of the specimen with a spacing

of 30 to 35 mm, depending on the specific specimen. Current was provided

through threaded end caps from a bipolar power supply with a 25 -A capability. A precision shunt was monitored with a conventional DVM to supply current readings. Voltage measurements were made with the current

flow in the sequence forward—reverse—forward. The program monitored

drift of the readings and remeasured if the variation between the two

forward readings exceeded 0.1 percent. In initial tests, a series of

current values was used at each temperature to determine when heating of

the specimen first occurred. These tests led to use of a specimen cur- rent of 2.0 A at room temperature and 10.0 A at lower temperatures ; this

gave voltages in the microvolt range. Owing to nonuniformities in the

specimen cross section, uncertainty of the data is ±1 percent.

Cold Rolling

Tensile specimens, in increments of cold rolling from 10 to 60 percent, were obtained for CDA 104. The alloy was received as rolled plate

(32 mm thick, 76 mm wide, 280 mm long) in the half -hard condition; then

it was annealed at 500°C for 90 min and air-cooled. The annealed plate

was uniaxial ly cold-rolled to a 10 percent reduction in thickness; this was achieved in approximately five passes through 5000 -kg, 150-mmdiameter rolls. Material for four tensile specimens was then removed

148
from the plate. The plate was again incrementally and uniaxially coldrolled another 10 percent, and material for four more tensile specimens

was removed. This procedure was repeated to obtain 30- , 40- , 50- , and

60-percent cold-rolled conditions.

Grain Size Control

The as-received, half-hard condition, plate material was uniaxially

cold-rolled to 50-percent reduction in thickness. Tensile specimens were

machined from the cold-rolled plate and recrystallized to selected grain

sizes in a salt bath (sodium nitrite and potassium nitrate). Recrystallization of the cold-rolled plate occurred at temperatures as low as 410°C , but at the lower temperatures there was very little grain growth.

The recrystallization and grain- growth time-temperature schedules are summarized in Table 1.

Table 1. Recrystallization Schedule and Grain Sizes

Number

of

Specimens

Temperature

( 8 C)

Time

(min) Cooling

Resulting

Grain Size

(Atm)

4 625 1 Air 23

4 625 5 Air 33

4 625 30 Air 44

4 720 90 Air 102

RESULTS AND DISCUSSION

Cold Rolling

The effects of cold rolling on tensile yield strengths at 295, 76,

and 4 K are shown in Fig. 1. In the annealed condition, 10-percent cold

rolling increased the yield strength substantially—by almost one order

of magnitude. Only moderate temperature dependence was exhibited. The

primary contributor appears to be the long-range stress fields (cell

structure) that impede dislocation motion. These long-range stress

fields are firmly established at 10 percent strain; subsequent deformation does not contribute as significantly to yield-strength increases.

The initial rapid increase of yield strength with cold rolling, followed

by a leveling off at higher deformations, has also been reported by

Truckner and Mikkola3 and Zehetbauer and Trattner 4 for oxygen-free, highconductivity copper at room temperature. These studies included even

larger deformations and reported eventual softening at strains greater

than about 1.5.

149
500

Fig. 1. Effects of cold-rolling CDA 104 copper

on tensile yield strength.

The effect of cold rolling on tensile ultimate strength is illustrated in Fig. 2. In contrast with the effect on yield strength,

(1) ultimate strength does not increase abruptly at small strains and

(2) its dependence on cold work is lowest at low temperatures. Presumbly, the difference in dependence at small strains reflects the

parabolic nature of copper stress—strain behavior. The stronger dependence of ultimate strength on deformation at higher temperatures probably

reflects larger recovery effects at higher temperatures—during tensile

deformation as opposed to during cold rolling.

Cold rolling of copper does not significantly affect tensile reduction of area at low temperatures (Fig. 3). There is virtually no dependence on temperature of reduction of area, especially at larger

amounts of cold work. Since reduction of area normally correlates with

fracture toughness for ductile, face-centered cubic alloys, the implications of these data are that toughness of copper is not affected by cold

work. However, we must caution that there is strong correlation between

yield strength and fracture toughness; therefore expect a decrease of

fracture toughness with cold work on those grounds.

150
Fig. 2. Effects of cold-rolling CDA 104 copper

on tensile ultimate strength.

Fig. 3. Effects of cold-rolling CDA 104 copper on tensile elongation and reduction of area.

151
Tensile elongation of copper is inversely related to temperature

(Fig. 3). Not only is the elongation highest (per unit cold work) at

4 K, but this increased ductility is generally maintained at larger increments of cold work. There is stronger dependence of elongation on cold work at lower temperatures, similar to that of the yield strength

but in contrast to that of the ultimate strength.

The electrical resistivity data in Table 2 appear to be consistent

with the behavior of the oxygen- free copper that we measured earlier. 5 This was expected, since the two coppers are quite similar in composition. In the earlier experiment, the shape of the RRR—deformation curves differed significantly for copper material from different sources. The

4-K data represent the contribution from the dislocations introduced

during cold rolling. The Ap for the 10- and 20-percent cold-rolled con- ditions are greatest, corresponding to the more rapid increase of yield

strength during these increments.

Table 2. Resistivity of Cold-Rolled CDA 104

Reduction

(%) 295 K

Resistivity

76 K

(/ifl-cm)

4 K

0 1.75 0.206 0.0100

10 1.74 0.208 0.0168

20 1.76 0.219 0.0253

30 1.74 0.221 0.0279

40 1.75 0.223 0.0295

50 1.74 0.226 0.0320

60 1.76 0.231 0.0355

Grain Size

The yield and flow strengths of most metals and alloys increase with

decreasing grain size, following the well-known Hall—Petch relation:

<7y - aQ + kd”^ (1)

where aQ

is the polycrystalline strength extrapolated to infinite grain

size, d is the average grain diameter, and k is the dependence of

strength on d*

3*. Both <7y and k usually depend on strain and temperature.

This dependence is shown in Fig. 4 for CDA 104 copper at 295, 76, and

4 K. The slope, k, increases at lower temperatures; Hansen and Ralph6 reported a lower slope at small strains (less than 0.05) at 76 K than at

room temperature, but a higher slope at larger strains. Our Hall—Petch

terms are summarized in Table 3; also included this in table are the re- sults of a data review of CDA alloys 101—104 by Simon and Reed7 and the

152
Fig. 4. Effects of grain size on the tensile

yield strength of CDA 104 copper.

Table 3. Hall-Petch Parameters from Various Studies

Source Temperature

(K)

Parameters

oq (MPa) k (MPa/Mm* 4

)

This paper 295 7.5 142

Simon, Reed7 295 18.6 112

Hansen, Ralph6 *

295 20 152

Merz, Dahlgren8 295 6.4 149

This paper 76 11 180

Hansen, Ralph * 76 -45 -

This paper

. ;—:—:—= 4 11 188

^Extrapolated from strains >0.03

153
data of Merz and Dahlgren8 who studied extremely fine grain sizes. Simon

and Reed conducted regression analysis of all available data at room

temperature using the Hall—Petch equation.

From Table 3 we can conclude that the slope of the Hall—Petch relationship between yield strength and grain size is fairly well-established

at 295 K; it ranges from 142 to 152 MPa/pm^ in the three studies. At low

temperatures k increases, reaching about 188 MPa//im‘^ at 4 K. The con- stant, o0 , depends on specimen purity. Our data from specimens of con- stant purity indicate a modest increase of o0 from 295 to 76 K, but

little change of aQ between 76 and 4 K.

Grain size only modestly affects ultimate strength; as shown in

Fig. 5, this dependence follows a Hall—Petch relationship. The increase

of oQ with decreasing temperature reflects the stronger temperature dependence of the ultimate strength. The slope of the dependence is approximately independent of temperature. This is unexpected, since k

at much lower strains is dependent on temperature (Table 3) , and Hansen

and Ralph reported a strong dependence of k on strain at 76 K. The

Fig. 5. Effects of grain size on the tensile

ultimate strength of CDA 104 copper. 5

154
implications are that grain boundaries and strength level play a strong

role in the ductile fracture of copper tensile specimens and that this

role is independent of low temperatures.

Tensile ductility, as measured by elongation and reduction of area,

is independent of grain size in most cases, and reduction of area is

practically independent of temperature (Fig. 6). The only apparent dependence is the effect of grain size on tensile ductility at 4 K: decreasing grain size reduces elongation (while increasing the tensile

strength), and reduction of area remains constant. Deformation twinning,

probably responsible for the serrated yielding at 4 K, may be responsible

for this trend.

The electrical resistivities of copper with various grain sizes are plotted in Fig. 7. The 102-/im grain-size specimen, represented by the

solid circle, was annealed at 720°C and then annealed at 25°C to give the

circled point for that grain size. Error bars on the nitrogen and helium

graphs represent measurement precision. On the 295-K graph, it represents the error possibly due to a ±0.5-K temperature change.

The resistivity measurements show a clear, if small, dependence of

that parameter on the measured grain size for this copper. Two interesting features of Fig. 7 are the larger dependence of the resistivity on grain size at 295 K than at the lower temperatures and the relatively

Fig. 6. Effects of grain size on the tensile reduction of

area and elongation of CDA 104 copper strength.

155
Fig. 7.

Eo

<3>

H

>

CO

COw

CE

1.78 1 1 1 1 1 CDA 104

1.76 o Room Temperature

I 1.74

o 8

1.72

o

t

* T

0.210 -

Liquid Nitrogen

i 0.205 _ O _

° o

o 0.200

0.020

* ^

•

Liquid Helium

0.015

o

M d o o

0.010

l _l 1 1 1—

0 20 40 60 80 100 120

GRAIN SIZE.pn

Resistivity of CDA 104 copper as a function of grain size.

large change in the resistivity caused by annealing the specimen with the

largest grain size at a higher temperature (720°C) than that used for the

other specimens (625°C). In the case of the 295-K resistivity, one is

tempted to postulate a temperature -dependent grain-boundary resistivity

or an annealing effect on the lattice resistivity. Neither of these con- cepts is supported by previous work, to the best of our knowledge, and

our data are too few to enable a detailed evaluation. The rather large

increase in the residual resistivity observed to result from the higher

temperature anneal of the largest grain- size specimen indicates that some sort of precipitation or strong clustering of resistive impurities is

taking place. Reannealing this specimen at the lower temperature apparently removes the impurity, perhaps by reformation of an agglomerate.

The resistivity of the smaller grain- size specimen is above the expected

value for the fully annealed state (1.72—1.74 /xG- cm) . One explanation

for this higher resistivity is that the initial cold- worked specimen was not completely recrystallized; however, the decrease of the lower temperature resistivities for the smallest grain size specimen do not

support this interpretation.

156
High -Conductivity Alloy Strengthening

Usually, the objective in strengthening of high- conductivity metals

is to maximize strength while maintaining minimum resistivity. Consider

the ratio of yield- strength increase (Aay) to resistivity increase (Ap)

at low temperatures. Attendant with grain size decreases from 102 to

23 /im, the ratio Aay/Ap is about 5,400 (MPa//ifl-cm) at 76 K and 19,600 at

4 K. After cold rolling, this same ratio ranges from about 37,000 at 4 K

for 10-percent reduction in thickness to 15,600 at 76 K for 60-percent

reduction in thickness. For low- temperature use, 10 -percent cold rolling

appears to optimal. Since both strengthening contributors produce nearly

equivalent strength—resistivity ratios, it is likely that both grain- size

and cold- rolling effects are desirable in terms of dislocation contributions. For 295 -K applications, cold rolling appears to be superior,

since neither reduction of grain size nor cold rolling contribute substantially to increases in resistivity.

CONCLUSIONS

From tensile and electrical resistivity measurements at 295, 76, and

4 K of 99.95 copper with varying amounts of cold rolling and differing

grain sizes, the following conclusions emerge:

1. Cold -rolling dramatically increases yield strength, ultimate

strength and electrical resistivity, reduces elongation, but does

not affect reduction of area. Cold rolling to 10-percent reduction

in thickness increases the yield strength by a factor of about 8 at

all low temperatures; subsequent rolling to 60 percent increases the

yield strength by a factor of only about 1.5.

2. Yield and ultimate strengths are related to d'*1

, where d is the

grain size. This dependence increases at low temperatures. Ductility is unaffected by grain size, except for elongation at 4 K.

Grain size has little effect on resistivity.

3. On the basis of optimal strength and minimal resistivity, a modest

amount of cold rolling (about 10 percent) produces the best ratio of

strength- to-resistivity at low temperature.

ACKNOWLEDGMENTS

Many thanks are due to L. Scull who assisted in cold-rolling the

copper plate and K. Yu for electrical resistivity measurements. This

study was partially supported by U.S. Department of Energy and the International Copper Research Association.

157
REFERENCES

1. R. P. Reed and R. P. Walsh, Low temperature tensile characteristics

of copper and aluminum in "Magnetic Studies for Magnetic Fusion

Energy Applications at Low Temperatures — IX," NBSIR 86-3050,

National Bureau of Standards, Boulder, Colorado (1986) pp. 107—140.

2. F. R. Fickett, Electrical properties, chapter 5 in "Materials at Low

Temperatures," R. P. Reed and A. F. Clark, eds. American Society for

Metals, Metals Park, Ohio (1983), pp. 161—201.

3. W. G. Truckner and D. E. Mikkola, Strengthening of copper by dislocation substructures, Metall. Trans. 8A, 45—49 (1977).

4. M. Zehetbauer and D. Trattner, Strengthening characteristics of

heavily cold-rolled copper, in "Strength of Metals and Alloys,"

Pergamon Press, New York (1985), pp. 105—110.

5. F. R. Fickett, The effect of mill temper on the mechanical and

magnetoresistive properties of oxygen- free copper in liquid helium,

Adv. Cryog. Eng. — Mater. 30, 453—460 (1984).

6. N. Hansen and B. Ralph, The strain and grain size dependence of the

flow stress of copper, Acta. Metall. 30, 411—417 (1982).

7. N. J. Simon and R. P. Reed, Properties of copper and copper alloys,

to be published, National Bureau of Standards, Boulder, Colorado.

8. M. D. Merz and S. D. Dahlgren, Tensile strength and work hardening of

ultrafine- grained high-purity copper, J. Appl. Phys. 46, 3235-3237

(1975).

158
TENSILE AND CREEP-FATIGUE PROPERTIES

OF A COPPER-STAINLESS -STEEL LAMINATE

L. L. Scull and R. P. Reed

National Bureau of Standards

Boulder, Colorado

The design of compact- ignition magnets uses a highconductivity copper alloy. Since the large magnetic fields

cause large stresses in the coil, this application may require

a conductor with higher strength than that of the copper alloys

and equally high electrical and thermal conductivity.

A candidate material was produced by reinforcing the

copper alloy with a stainless-steel alloy. The steel is rollbonded into the midplane between two copper sheets. This material has the high thermal and electrical conductivity of the

copper alloy and, possibly, sufficient strength to be used in

compact -ignition magnets.

Tests were conducted at 295 and 76 K to characterize the

tensile and creep—fatigue behavior of this laminated composice

material in three roll-bonded conditions. These conditions

correspond to a bulk reduction of 40, 50, and 60% cold work

in the laminate as a whole. A mixing rule was used to predict

the tensile behavior of the composite on the basis of the individual tensile properties of copper and stainless steel.

INTRODUCTION

Compact- ignition magnets are produced by winding a helical conductor

into a torus. When pulsed, the magnet is subjected to large thermal

and mechanical stresses. The large current density causes a substantial

amount of Joule heating within the conductor, and the magnetic field

produces a large hoop stress within the torus. Copper alloys possess the

required electrical and thermal conductivity, but they may not have sufficient strength for this application.

In Advances in Cryogenic Engineering — Materials

,

vol. 34, Plenum

New York, 1988, pp. 397-403.

159
A candidate material with both high strength and conductivity has

been produced by roll-bonding a stainless steel alloy between copper

sheets. The magnet conductors are cooled with liquid nitrogen (76 K) to

reduce the electrical resistance and Joule heating after the magnet is

pulsed.

Tensile tests were conducted to determine the strength of this com- posite at 295 and 76 K. The creep—fatigue behavior was also investigated

to determine the effects of cyclic loading caused by the magnetic pulses.

SPECIMEN CHARACTERIZATION

A stainless steel alloy (Fe—18Cr—13Mn-3Ni—0 . 37N) was roll -bonded

into the midplane between two copper sheets (99.95% Cu + Ag minimum).

Three types of specimens were tested in conditions corresponding to bulk

thickness reductions of 40, 50, and 60% in the roll-bonding process.

Tensile tests were conducted on specimens with a uniform width in the

gauge length and creep—fatigue tests were conducted on specimens with a uniform radius in the gauge length (see Fig. 1).

I cm

Figure 1

.

Specimen geometry of laminates tested in this study.
TEST APPARATUS AND PROCEDURE

Uniaxial tensile tests 1 were conducted at a constant displacement

rate of 0.3 cm -min* 1

, which corresponds to a strain rate of 5 x 10* 2 min* 1 at 295 and 76 K. The specimens were installed in conventional

wedge grips on a servohydraulic load frame for the 295-K tests. An ex- tensometer with a 38 -mm (1.5 -in) gauge length was used on all tensile

tests

.

The 76-K tests were conducted using a screw-driven load frame, which

was fitted with a cryostat to enable the test specimen to be completely

immersed in liquid nitrogen. The 76-K specimens were pinned and clamped

in specially constructed grips designed for low- temperature applications.

The creep—fatigue tests were conducted on a servohydraulic load

frame using the same grips as for the 295-K tests. All creep—fatigue

tests were conducted at room temperature owing to expected specimen- to

specimen variations and a limited specimen supply. This gave two replicates for each cold-worked condition tested.

A waveform generator was used to program the hydraulic system in a load-controlled servo loop. The specimen stress was ramped up from 9 to 90% of the yield stress of the composite being tested (40, 50, or 60%

reduction) in a period of 5 s. The 90% stress was held constant for

10 s, and then ramped down from 90 to 9% of the yield stress in 5 s . The

stress was—then held constant at 9% for an additional 10 s. This loading

process was repeated several thousand times.

A strain gauge mounted at the center of the specimen across the

thinnest portion of the gauge length measured the strain in the specimen

as a function of the number of loading cycles. The same extensometer

used in the tensile tests was also used as a less accurate, but independent, measure of the specimen strain.

RESULTS

The measured and predicted values of Young's modulus (E) and the

0.2% offset yield strength (<7y) at 295 K are given in Table 1. The

predictions were made by using a simple mixing rule that states: the

strength (modulus) of the laminate equals the sum of the strengths

(moduli) of each layer multiplied by the fraction that each layer forms

in the laminate. Table 2 contains similar data obtained at 76 K. The

parameter a is the fraction of the thickness of the laminate occupied by

the copper.

Stress—strain curves for the laminate at 295 and 76 K are shown

in Figs. 2 and 3, respectively. These curves indicate that the strength

of the laminate increases with increasing cold work and with decreasing

temperature. Figure 2 shows small spikes near failures at stresses of

approximately 0.2 GPa. These spikes occurred when the copper arrested

161
Table 1. Predicted and Measured Mechanical Properties at 295 K.

Cold Work,

%

E (meas . )

,

GPa

E (pred. )

,

GPa

£7y (meas .

)

MPa

<7y (pred. )

,

MPa

a

40.1 106 144 526 582 0.7214

49.0 144 142 470 650 0.7418

60.3 137 142 593 663 0.7417

Table 2 . Predicted and Measured Mechanical Properties at 76 K.

Cold Work,

%

E (meas . )

,

GPa

E (pred.),

GPa

<7y (meas .

)

MPa

<7y (pred. )

,

MPa

a

40.1 113 151 620 801 0.7214

49.0 101 149 618 863 0.7418

60.3 114 149 604 879 0.7417

Figure 2. Stress—strain curves at 295 K.

162
the fracture in the tensile bar immediately after the stainless steel

failed. The copper then was pulled to a knife edge with some strain

hardening prior to final failure.

Figures 4 through 7 show Young's modulus, 0.2% -offset yield

strength, ultimate tensile strength, and elongation to fracture, re- spectively, as a function of bulk cold work in the specimens tested.

Figure 8 shows the curves of strain versus number of cycles that were

obtained in the creep—fatigue tests.

Figure 3. Stress—strain curves at 76 K.

Figure 4. Young's modulus versus cold work at 295 and 76 K.

163
ULTIMATE

STRENGTH,

MPa

H

YIELD

STRENGTH

(0.2%),

MPa

5. Yield strength (at 0.2% offset) versus cold work

at 295 and 76 K,

Ultimate tensile strength versus cold work

at 295 and 76 K.

1 64

Figure 6.
DISCUSSION

A mixing rule was proposed to enable predictions of laminate

strength as a function of the strengths of the layers and their respective fractions of the thickness of the laminate. 2 This rule can also be

applied to the modulus of elasticity. The mixing rule appears to over- estimate the strengths measured in the tensile tests, as seen in Tables 1 and 2. The modulus data3

" 4 and yield strength data5

" 6 were obtained by

extrapolating existing data.

The modulus was insensitive to cold work, as shown in Fig. 4, but it

increased slightly as the temperature was decreased, as would be

expected. The yield strength increased slightly with increased cold work

at 295 K only (Fig. 5). This increase was not seen at 76 K, probably

owing to thermally induced stresses caused by the difference in thermal

expansion coefficients between the copper and stainless steel. Ultimate

strength increased significantly with increasing cold work and decreasing

temperature (Fig. 6). Elongation to fracture decreased slightly at 295 K

with increasing cold work, as expected (Fig. 7). However, the elongation

increased with increasing strength at 76 K. This may have been caused by

the difference in the thermal expansivities, which affect the residual

stress state in the laminate upon cooling.

Creep—fatigue data for the average of two tests for each coldworked condition are shown in Fig. 8. More creep was observed as the

cold work was increased, which was not expected. However, the fraction

of stainless -steel thickness decreased as the cold work was increased

(see Tables 1 and 2). Thus, with increased cold work, there was less

stainless steel to support the loads, and more creep occurred. All

specimens seemed to reach a degree of creep saturation after a large

number of fatigue cycles.

CONCLUSIONS

In a copper—stainless -steel laminate, minor variations were observed

in Young's modulus; these variations were probably due to specimen bending caused during the roll-bonding process. The yield strength increased

slightly from 295 to 76 K at the lower cold-worked conditions. The ultimate tensile strength increased substantially from 295 to 76 K. At 76 K,

fracture occurred at the ultimate strength. At 295 K, the fracture was arrested by the copper. The fatigue values were somewhat scattered,

probably as a result of local bonding variations in the laminate . Creep

seemed to saturate after a large number of fatigue cycles. The mixing

rule provided a method to estimate composite behavior. However, the predictions were not conservative, so adjustments are required to improve

the method.

165
Figure 7. Total elongation to fracture versus cold work at 295 and 76 K.

I 5 10 50 100 500 1000 5000

CYCLES

Figure 8. Creep—fatigue behavior at 40, 50, and 50% cold work at 295 K.

166
ACKNOWLEDGMENTS

This research was sponsored by the Office of Fusion Energy, U.S.

Department of Energy. The authors are indebted to Dr. H. Becker, Massachusetts Institute of Technology, for the helpful discussions and for

supplying the test specimens.

REFERENCES

1. Standard Methods of Tension Testing of Metallic Materials, ASTM E8-

82, American Society of Testing and Materials, Philadelphia, 119

(1984).

2. H. Becker, M. Besen, E. Bobrov, J. Chen, and D. B. Montgomery,

Structural Properties of Reinforced Copper, PFC/RR- 86 - 12 , Massachusetts Institute of Technology, Cambridge, Massachusetts (1986).

3. H. M. Ledbetter, Elastic constants of polycrystalline copper at low

temperatures, Phvs . Status Solidi 66:477—484 (1981).

4. H. M. Ledbetter and E. W. Collings, Elastic constant anomalies at the

Neel transition in Fe—18Cr-3Ni-12Mn, Mater. Sci. Eng. 68:233-237

(1984-1985).

5. R. P.HReed, R. P. Walsh, and F. R. Fickett, Effects of Grain Size and

Cold Rolling on the Cryogenic Properties of Copper, in "Advances in

Cryogenic Engineering - Materials," vol. 34, Plenum, New York, 299-

308 (1988).

6. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and Tensile Properties of Three Nitrogen- strengthened Stainless Steels at

Cryogenic Temperatures," NBSIR 78-884, National Bureau of Standards,

Boulder, Colorado, 91-154 (1978).

167
»

" -
WELDING

169

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, a yield strength

of 991 MPa and a fracture toughness of 287 MPa-m^ were measured for a gas

metal arc weld, substantially higher than the 203 MPa-m^ reported last

year. The exceptionally high toughness in the Fe—18Cr—20Ni—5Mn—0 . 16N

weld metal was achieved with a new welding procedure that decreases the

oxygen content of the weld, in this case, to one-eighth that of a weld

by other procedures. The lower oxygen content is related to smaller and

fewer inclusions.

• A diagram was developed to predict the ferrite number (FN) in stainless

steel welds. When it was evaluated with a database containing almost a thousand compositions, the diagram was substantially more accurate than

the DeLong or Schaeffler diagrams. It is applicable from 0 to 100 FN and

up to 10 wt.% manganese, 3 wt.% molybdenum, and 0.2 wt.% nitrogen. The

new diagram is particularly useful for stainless steel welds in which

fissuring resistance and fracture toughness must be controlled by the FN

• Electron beam and laser welds were evaluated in two 25 -mm- thick copper

alloys, C10700 (Cu—0.08Ag) and C17510 (Cu—0 . 4Be—1 . 8Ni) , candidate materials for use in the Compact Ignition Tokamak. A substantial drop in the

yield strength of transverse tensile specimens occurred at both test

temperatures, 76 and 295 K. By monitoring a series of 1-mm-wide regions

in the gage length, we determined that the lower strength after welding

was due to initial yielding in the very soft weld and HAZ regions.

• Yield and notched- fatigue strength were measured for stainless steel

tube- to -flange welds at 4 and 295 K. Three flange materials, AISI 304L.

AISI 316L, and 21Cr—6Ni—9Mn, were joined to 21Cr—6Ni—9Mn tube by gas

tungsten arc welding. Welds with the 316L flange material had the be-.

yield and fatigue strengths, close to that of the tube material, at bo-h

temperatures

.

171
• An aluminum- deoxidized weld was compared with conventional silicondeoxidized welds to clarify the role of inclusions in fracture. The

inclusions in the aluminum- deoxidized weld were more resistant to

void formation, but the poor strength—toughness combination (900 MPa85 MPa-m^) of this weld at 4 K is attributed to the large amount of

residual aluminum in the matrix.

172
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS

Weld Alloy Welding Tensile Charpy Fracture Fatigue Crack

* Report Volume

[designation] Process Impact Toughness Growth Rate (page)

AUSTENITIC STAINLESS STEELS

Fe-13Cr-20Ni-9Mn-2Nb SMAW I, II I, II, I (159,169)

VIII II (299) VIII (29)

Fe-1 5Cr-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 IV (275) 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-1 7Cr-16Ni-4Mo SMAW IV IV rv (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 m.itf IV III, IV III, VI III (155)

IV (275,453)

V (199)

Fe-18Cr-18Ni-2Mo-3Mn-N SMAW IV IV IV (453)

Fe-18Cr-20Ni-5Mn-N GMAW IX IX IX (247)

Fe-18Cr-20Ni-6Mn-0.3Nb GMAW IV IV rv 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 (275) VIII (29)

GTAW 111,1V III, IV, IV III (155.195)

VIII IV (275) VIII (29)

SAW III, IV, III, IV, III, IV III, VI III ( 155,195) V V IV ( 275.415)

V ( 233)

VI ( 199)

^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)

.

RBSIR 82-1667 (V) ; NBSIR 83-1690 (VI); NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII).

NBSIR 86-3050 (IX); NBSIR 87-3067 (X).

173
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS, continued

Weld Alloy

[designation]

Welding

Process Tensile CharpyImpact Fracture Toughness

Fatigue Crack

Growth Rate

k Report Volume

(page)

Fe-19Cr-12Ni-2Mo SMAW II, III, II, VIII II, III, III, IV, VI II (299,315)

[316L], continued IV IV III (155,167)

VI (199)

* VIII (29)

Fe-19Cr-12Ni-2Mo-N SMAW II II, VIII II II (299)

[316LN] VIII (29)

Fe-19Cr-16Ni-7Mn GMAW IV IV IV (453)

Fe-20Cr-10Ni SMAW I I, VIII I (169)

[308] VIII (29)

Fe-20Cr-10Ni FCAW III III III, VI III (155)

[308L] VI (199)

GMAW III III III, VI III (155) VI (199)

SMAW II, III II, VIII II, III III, VI II (299)

III (155)

VI (199) VIII (29)

Fe-20Cr-10Ni-N FCAW IV IV IV, VI IV (275,415)

[ 308LN] V (199)

Fe-20Cr-10Ni-9Mn-N SMAW I I I (159,169)

Fe-20Cr-16Ni-7Mo-2W SAW IV IV IV (275)

Fe-20Cr-34Ni-2Mo-Nb SMAW VII VII VII (253)

Fe-21Cr-6Ni-9Mn-N SMAW I, II I, II I (159,169)

II (299)

AUSTENITIC STEELS

Fe-1 5Mn-8Ni-lMo-0 . 7C SAW VIII VIII VIII (15)

Fe-5Cr-25Mn-l|i 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 IV III (155,217)

[5183] IV (323)

AI-5Mg-0.12Ti GMAW III IV IV III (155, 217)

[5556] IV (323)

*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).

174
FERRITE NUMBER PREDICTION TO 100 FN

IN STAINLESS STEEL WELDS*

T. A. Siewert and C. N. McCowan

National Bureau of Standards

Boulder, Colorado

D. L. Olson

Colorado School of Mines

Golden, Colorado

To improve the accuracy of ferrite number (FN) prediction

in stainless steel welds, a new diagram has been developed

from a database containing more than 950 alloy compositions

from worldwide sources. In accuracy, this diagram surpasses

the DeLong diagram for the low-FN austenitic-stainless-steel

compositions of the AISI 300 series, and it corrects a 2-FN

bias detected for AISI type 309 stainless steel. The diagram is more accurate than the Schaeffler diagram for duplex

stainless-steel alloys and ferrite contents to 100 FN. It is

suitable for Mn contents to i0 wt.%, Mo contents to 3 wt . %

,

N contents to 0.2 wt.%, and Si contents to 1 wt.%. Changes in

the primary solidification mode are indicated on the diagram;

they appear to affect the FN response. Transitions in iso-FN

line spacings may be caused by these mode changes.

I . INTRODUCTION

In many applications, the ability to control the delta- ferrite

content of stainless steel welds is important. For example, the ferrite

number (FN) is often used as an indicator of the resistance to fissuring

(hot tearing) that occurs in many of the AISI 300-series stainless

steels. 1 Fissuring decreases rapidly over a small range of ferrite

contents in these primarily austenitic phase weld compositions, so an accurate predictor of the FN is necessary to ensure the minimum ferrite

levels that will prevent cracking. Fissuring should not be avoided by

k Submitted to Welding Journal.

175
simply using materials with high FN values, since higher FN levels con- tribute to other undesirable behavior, such as an increased corrosion

rate in certain environments. 2 Thus, there is incentive for keeping the

FN in this limited range and for controlling the ferrite accurately.

At cryogenic temperatures, ferrite reduces the Charpy V-notch absorbed energy and the fracture toughness of stainless steels. 3

’ 6 The

response is as dependent on ferrite content as is fissuring, but it

reacts in the opposite direction. To provide the appropriate balance

of cracking resistance and fracture toughness, the ferrite content must

be controlled in a narrow range. Cryogenic applications require the

most stringent FN controls and justify a search for more accurate FN

prediction.

These examples have all been in the 0 to 25 FN range. In the past

ten years, we have seen the development of new duplex stainless steel

compositions for service conditions under which the austenitic grades

are unsuitable. As the name suggests, these alloys contain substantial

amounts of both the austenite and ferrite phases, with the ferrite con- tent being as high as 100 FN. For these alloys, Kotecki has recently

shown relationships between the ferrite content and various mechanical

properties, such as tensile strength and elongation. 7 So, in the high

FN range also, it is important to control the ferrite content to achieve

optimal weld properties. Unfortunately, fewer predictive diagrams have

been developed for these compositions, and the data with which they were developed may not be accurate for today's compositions.

In the United States, the DeLong8 and Schaeffler 9 diagrams are probably the most common methods for predicting the ferrite content of

stainless steel welds. For both diagrams, a position on the diagram

is determined from the composition and then related to a series of isoferrite or iso-FN lines. Both diagrams list the elements promoting the

austenite phase along the vertical axis (as a term called Ni e q) and the

elements promoting the ferrite phase along the horizontal axis (as a term

called Cre q) . This nomenclature is based on naming the term after its

primary component (Ni or Cr) and adding the other components, with appropriate coefficients to represent their relative power in promoting aus- tenite or ferrite, to form an equivalent (Ni e q or Cre q) . The Schaeffler

diagram (Figure 1) uses straight lines that pivot around a common focus

to represent the amount of ferrite as volume percent over the range 0 to

100%. 9 The most recent version of the DeLong diagram (Figure 2) uses parallel Tines to represent the amount of ferrite as FN over the range

0 to 18 FN. 8

Earlier versions of the DeLong diagram represented ferrite in units

of volume percent only. This recent version (Figure 2) was produced

after the Welding Research Council (WRC) Subcommittee on the Welding of

Stainless Steels promulgated a standard using the term ferrite number

(FN) to define the ferrite content by its magnetic response. As a re- sult, the DeLong diagram (now) contains both the FN and volume percent

ferrite designations, and it indicates that the two measurement systems

begin to diverge at ferrite contents near 8 FN.

176
NICKEL

equivalent

=

y.Ni +

30x

y.c +

30x

y.N

-f

o.5

x

%Mn

’

Nickel

Equi

volenl

•%

Nj

+

30X

%c+

0.5X

Mn

1. The Schaeffler constitution diagram for stainless - steel weld metal. 9

Fig. 2. The DeLong constitution diagram, revised January 1973 to convert

it to the Welding Research Council's ferrite-number svstera for weld metals

.

8

177
The two diagrams differ only slightly in their representation of

compositional effects on ferrite content. The Cr e q

expressions are identical in the two diagrams—even the coefficients are the same. The

DeLong Ni e q expression contains all the elements and coefficients of the

Schaeffler diagram but adds a term to include the effect of nitrogen

on FN.

The difference in the applicability of the two diagrams is substantial. The DeLong diagram is a finely tuned subset of the Schaeffler

range, designed specifically for the AISI 300-series stainless steels

containing small amounts of ferrite. Its widely accepted use with these

alloys is documented by its inclusion in the American Welding Society

filler metal specifications A5.4 and A5 . 9 and in the American Society of

Mechanical Engineers (ASME) Boiler and Pressure Vessel Code. Prediction

of FN for the compositions, such as the duplex stainless steels, that

fall outside the range of the DeLong diagram seems to be the most common

quantitative use of the Schaeffler diagram; however, its extension to

fully ferritic compositions and the inclusion of martensitic regions

provides much qualitative information about the general effect of alloying elements and the characteristic mechanical properties of alloy

families

.

Inaccuracies in both the DeLong and Schaeffler diagrams have become

evident with the generation of new alloys and data. As a result, revisions of these diagrams have been proposed and several alternate diagrams

have been developed. 10 The 0.5 coefficient for Mn is questioned by a number of researchers. Hull proposed a quadratic form indicating that

Mn promoted austenite less as its content was increased, and actually

promoted ferrite at very high levels. 11 Espy, in a study of stainless

steels with up to 12.5 wt.% Mn and 0.31 wt.% N, concluded that Mn should

not have a coefficient: its austenitizing potential could be better

represented as a constant. 12 In a study of 41 specially formulated

shielded-metal-arc electrodes with Mn levels gradually increasing to

13 wt.%, Szumachowski and Kotecki 13 also determined that the Mn aus- tenitizing potential could be represented as a constant, but suggested

a value slightly smaller than Espy's. From a matrix of 26 electrodes

containing up to 10 wt.% Mn and 0.26 wt.% N, McCowan and coworkers con- cluded that Mn and N interacted in a complex manner that could be ex- pressed best by Ni eq = Ni + 29(C + N) + 0.53(Mn) - 0.05(Mn) 2 - 2.37(MnN)

+ 0.94(MnN) 2 — 0.71. 14 However, they also indicated that the effect of

Mn could be adequately represented by a constant if a slightly poorer

accuracy was accepted.

Other studies attempted to understand the effect of N. Espy re- ported a variable N coefficient, decreasing in value from 30 to 20 as the N content increases from 0.2 to 0.35 wt.%. 12 Others reported singlevalued coefficients, ranging from 14.2 to 20. 11 * 15 ' 18 Ogawa and Koseki

reported a coefficient of 30 for FN but a coefficient of 18 for the

effect of N on solidification mode. 17

Recently Kotecki investigated the Si coefficient. 19 He found

that the coefficient of 1.5, indicated by both the DeLong and Schaeffler

178
diagrams, overstates the effect. In a study of six welds spanning 0.34

to 1.38 wt.% Si, he determined a coefficient of 0.1, which was significant for a 68% confidence interval. Another recent study found the Si

effect to be nonlinear; very small below 2 wt.% Si, but large at higher

Si contents. 20

Taken as a whole, these studies indicate that revisions of the

DeLong and Schaeffler diagrams can substantially improve the accuracy

of ferrite prediction for current stainless steel weld compositions.

Recent studies have emphasized that the solidification mode is more

accurate than the FN for predicting the resistance to hot cracking. 21-23

These studies emphasize that the initial solidification structure (primary austenite or ferrite) directly determines the sensitivity to cracking by controlling the solubility of elements that promote cracking at

these temperatures. These three studies contain data on both FN and the

primary solidification mode. Incorporation of these data into the database for this study enabled comparison of the two types of data.

The goals of this study were 1) to develop a database containing

recent FN data and new compositions, 2) to evaluate the DeLong and

Schaeffler diagrams with these data, 3) to determine which elements were

not being fit properly, and 4) to develop an improved predictive diagram

that was continuous over the range 0 to 100 FN (roughly equivalent to

65 vol.% ferrite). Ideally, this diagram would predict the ferrite con- tent of the conventional 300-series stainless steels with at least the same accuracy as the DeLong diagram and duplex steels with at least the same accuracy as the Schaeffler diagram, and it would be suitable for new

alloys containing extended compositional ranges. It would also include

primary solidification-mode data on the predictive diagram to reveal

relationships to the FN response.

II. DATABASE DEVELOPMENT

The approach to the development of the database was (1) to gather

a large sample of data for each of the common alloys that are now predictable by the DeLong diagram to ensure that these alloys would be

predicted accurately by the new diagram, and (2) to include data of many

special alloys to determine the limits of applicability and interactions.

Most of the data were obtained from members of the WRC Subcommittee

on Welding of Stainless Steel. The nine sources of data included electrode manufacturers and research institutes, as well as archival journals. The data well represented the common 300-series stainless steel

compositions: types 308, 308L, 309, 309L, 316, and 316L. These data

provided a good basis for evaluating the accuracy of the diagram and

comparing scatter within and among the various sources. Other data

represented duplex stainless steel compositions, experimental compositions designed to evaluate extended ranges for certain elements, and

other new alloys. These data were used to determine element interactions

179
and whether the behavior in the duplex range was a smooth extension of

the behavior in the low FN range

.

From the data sources, we requested an FN measurement by the procedure specified in AWS A4.2 24 or by the procedure developed by Kotecki 25

for extending the FN range to duplex alloys. The small fraction (approximately 1%) of the data for high-ferrite content alloys that was in

volume percent was converted to FN using Kotecki 's procedure. 25 Matching

each FN determination, we requested a complete chemical analysis specifically including C, Mn, Si, Cr, Ni, Mo, Ti, Nb , and N. These data have

been listed in the appendix to another report. 26 The compositional

ranges in this database are summarized by element in Table 1.

Table 1. Compositional Range for Various Elements

in the Database

.

Element Range (wt. %)

C 0.01 - 0.15

Mn 0.4 - 12

Si 0.1 - 1.3

Cr 15 - 32

Ni 5 - 25

Mo 0 - 7 N 0.03 - 0.9

Nb 0 - 0.9

Ti 0 - 0.1

Most of the data were received from manufacturers of welding electrodes, who had developed these data to certify that their products met

the customers' requirements. Since specifications controlled the analysis and measurement techniques and these specifications contained limits

for the critical parameters, the resulting data were quite consistent.

This precluded the variation in the FN of a given composition that David

et al. observed as the cooling rate was varied over several orders of

magnitude. 27 In fact, the great majority of data were produced with

shielded metal arc electrodes conforming to AWS specification A5 . 4 . Such

uniformity in the electrodes, welding process, welding procedure, and

180
measurement reduces variations in FN due to these variables and permits

a better fit of the data. Conversely, the predictive equation does not

consider these variables, and their effect must be determined before the

equation can be used to estimate the FN of welds produced under other

conditions

.

Although the procedure developed by Kotecki for measurement of FN 24 can be applied to fully ferritic alloys, they are outside the scope of

this study. This study was limited to about 100 FN since few data were

available above this level (almost none were supplied for the database)

,

and there does not seem to be a need to predict above 100 FN at present.

Some of the data included primary solidification mode information

following the nomenclature of Suutala2

1

(austenitic, austenitic—ferritic

,

ferritic—austenitic , and ferritic). These were projected on the FN predictive diagram for comparison of mode changes with iso-FN lines.

To evaluate the existing diagrams in the most careful manner and to

facilitate the location of errors, the database was divided into groups,

ten of which are listed below:

Group Description Cases

1. types 308 and 308L 220

2. types 309 and 309L 129

3. types 316 and 316L 139

4. FN < 18 796

5. FN < 18, and Mo < 3 wt . % 724

6. FN<18, N < 0.1 wt . %

,

and Mo < 3 wt.% 652

7. FN < 18, N < 0.1 wt . %

,

and Mn < 3 wt.% 662

8. FN < 18, N < 0.1 wt.%, Mn < 3 wt.%, and Mo < 3 wt.% 605

9. FN < 30, Nb > 0 443

10. FN > 18 126

By comparing the quality of the fit of these groups with the fit of

the entire database, we were able to determine the compositional range

over which the elements could vary without degrading the accuracy of the

predictive diagram. Similarly, these groups could be used to compare the

effect of the various elements in the new and old diagrams.

III. PROCEDURE

A variety of statistical representations were used to evaluate the

existing diagrams and the new expressions. To evaluate the existing

diagrams, relative-frequency histograms were used. The measured FN was subtracted from the FN predicted by the diagram, and the difference (AFN) was plotted versus the error class intervals (different levels of error)

A perfect fit would have resulted in a delta function centered at zero. The width of the distributions of the actual data around zero provided a measure of the variability, whereas an off-center peak revealed a bias.

181
The development of a new diagram began with a multiple linear re- gression, with FN as the dependent variable and the various elements that

affect the FN as independent variables. The process has been described

previously . 6 Briefly, it consists of the selection of the variable (element) that most strongly determines the FN and production of a predictive

equation. The program then chooses the next most influential variable

and develops a new equation. The process is repeated in a stepwise man- ner until the addition of further terms will not statistically improve

the prediction of FN. The coefficients for the various elements in these

equations were compared with the coefficients in the existing diagrams.

After evaluating the various groups of data individually, differences in

their coefficients could be used to determine which elements and compositions were the dominant sources of scatter in the overall predictive

diagram.

Next, the coefficient of each element was checked for accuracy.

Our approach was similar to a technique reported by Szumachowski and

Kotecki . 13 To evaluate the various coefficients that have been suggested

for Mn, they developed a term, ANie q, by subtracting from the measured FN

value the contributions of all other elements and plotting this term

versus the concentration of Mn. This ANie qterm was effectively the com- bination of the response of the FN to one particular element (Mn) con- founded with the random error in the data. When they were unable to

resolve a slope for the data, they concluded that Mn had no effect.

Our stepwise linear regression program produced a file of the re- siduals (error left after the fitting program was finished) for the data.

We plotted this versus the various elements that had not met the selection criteria of the regression program. For the elements that had been

selected by the regression program, we added the residuals to the product

of the element concentration and the selected coefficient. We called

this term ANie q

(or ACre qif the element functioned as a ferritizer)

since it contains the same information as the term used by Szumachowski

and Kotecki , 13 even though it was generated in a slightly different manner. The slope of the best- fit line through the ANie q or ACre q

versus

element concentration data was the coefficient for the element. This

technique allowed us to evaluate the quality of the fit and to search

for nonlinear behavior. By allowing various ranges of other elements, we also searched for element- interaction effects.

When we had incorporated the changes indicated by this technique

into our predictive equation, we used the equation to produce an FN map.

The map consisted of a blank FN diagram; its axes consisted of the Ni e q and Cre q from our best linear equation. The computer calculated a position for each datum from the Ni e q

and Cre qexpressions, and identified

that point by the measured FN. The program created a map by linear interpolation between measured FN values and then generated iso-FN contour

lines to represent the data. This technique enabled us to separate the

nonlinear response of the individual elements over the entire range of

the diagram from the nonlinear response in a specific region of the

182
diagram (i.e., to improve the fit by using either nonlinear coefficients

or curved and unevenly spaced iso-FN lines). The difference, AFN

,

between measured and predicted FN, was plotted as an error bar.

As a final check on the accuracy of the diagram, the AFN values

from the mapping program were plotted versus the element concentrations,

as was done after the linear fit by using the procedure for the ANi e q

and

ACre q

terms. By comparing these plots to the ANie q

and ACre qplots, we coula assure ourselves that the mapping procedure produced a diagram that

was at least of the same quality as the linear equation and determine

whether any nonlinear behavior could now be resolved.

The solidification-mode data were also plotted on the map to relate

the solidification mode to the FN behavior.

IV. RESULTS AND DISCUSSION

A. Construction of a New Diagram

To determine more accurate representations of the FN potential for

the type 308, type 309, and type 316 alloys, the multiple linear regression technique was applied to each data group. Three predictive equations were generated:

for 308 type alloys,

FN - - 25 + 2 . 8 (Cr

for 316 type alloys,

FN - — 15 + 2 . 2 (Cr

for 309 type alloys,

FN = - 40 + 3 . 6 (Cr

+ 0 . 9Mo) - 1 . 8 (Ni + 48C +30N)

;

+ 3Mo) - 2 (Ni + 40C + 30N)

;

+ 0 . 3Mo) - 2 (Ni + 55C + 23N)

.

(1)

( 2 )

(3)

The R2 values (a statistical measure of the fit that approaches 1 with

better fits) for these three equations were 0.65, 0.84, and 0.68; the

standard deviations of the residuals were 1.5, 1.2, and 1.6 FN respectively. The Cre q

and Ni e q

terms are the quantities in parentheses with

the various elements in units of weight percent. The differences in the

coefficients for the elements and the Cre q

and Ni e q

terms are due to differences in the behavior within the alloy range, as well as to difficulties in fitting a line to data that are in a small cluster. Because of

inaccuracies due to' data clusters, these equations are not recommended

for general use, but the equations do point out that each alloy range

generates a characteristic set of coefficients.

The differences in the equations also emphasize that an empirically

developed predictive diagram developed for many alloys can only approximate the behavior of a specific alloy group. Until we can develop a

183
model that truly represents the physical phenomena of the ferrite transformations, we must use a variety of fitting constants and statistical

techniques to develop a diagram that provides predictions of acceptable

accuracy. Although the optimum coefficients for each alloy group are different, the general equation forms are quite similar. This suggests

that a single equation with compromise values for the coefficients might

be able to predict the FN with an acceptable degree of accuracy.

Next, we processed the data of group 8 (FN'< 18, N < 0.1 wt.%,

Mn < 3 wt.%, and Mo < 3 wt.%) through the multiple linear regression

program. The group 8 data cover a broad range of compositions and produced equivalents with a range of applicability equal to that of the

DeLong diagram. From group 8, we excluded the data for FN < 1 and the

austenite—ferrite (AF) solidification mode, since we anticipated nonlinearities due to ferrite formed mechanically during surface preparation and

in differing solidification modes. The predictive equation, for group 8 alloys and alloys with FN > 1 , and solidifying in the ferrite—austenite

(FA) mode,

FN = - 18 + 2 . 9 (Cr + Mo + 0.3Si) - 2.6(Ni + 35C + 20N + 0.3Mn), (4)

is analogous to the DeLong diagram. For this equation, R

2 is 0.78, and

the standard deviation of the residual is 1.6 FN. These parameters are similar to those for the alloy groups considered above and suggest that

one equation can provide a fairly accurate prediction over a range of

compositions. This equation has coefficients similar to those in equations 1 through 3; in fact, it is almost the average of the behavior for

the three alloy groups. In equation 4, the coefficients of C and N within the parentheses are smaller, but when multiplied by the coefficient

outside the Ni e q

parentheses, the values are close to those of equations

1 through 3. The main difference is the larger effective coefficient for

Ni, 2.6 as compared to 1.8 or 2. Equations 1 through 4 all indicate that

the effect of C is greater than that of N, unlike the equal effects predicted by DeLong. The C/N coefficient ratios are 1.6, 1.3, 2.4, and 1.8

for these equations, respectively. Similar C/N coefficient ratios have

been found in other FN studies; a ratio of 1.5 is used in the Hammar

solidification equivalents

.

1

0

> 1

2

> 1

4

> 1

5

f Equation 4 is the best fit for a linear regression model, that is, an arithmetic combination of linear terms producing a diagram with equispaced and parallel lines. The following subsections investigate the

possibility of improving the prediction when thes'e requirements are relaxed.

1 . Changes in Nickel and Chromium Equivalents

Compositional terms defined by the best- fit linear equation (equation 4) were evaluated by using ANi e q

and ACre q

plots described in the

procedure. Although the technique was intended primarily to find secondorder or other nonlinear behavior, it also confirmed that the coefficient

determined during the linear regression program was the best line through

the data. We did note the existence of some interactions. Comparison of

184
ACr e q

for Cr and Mo plots revealed that Cr could be given a coefficient

of 1 and possess little scatter when the Mo content was less than 3 wt.%.

For higher Mo contents, the behavior was substantially different, with

a level of scatter that precluded determination of a mathematical

representation.

We were unable to confirm the nonlinear trends reported for Mn

(Hull 11

) or Mn-N (McCowan et al. 14

), but unless these trends were of

substantial magnitude, their effects would have been hidden by the scatter in the data collected from many sources. The ANi e q—Mn plots did indicate that scatter about the 0.3 slope expressed for Mn in equation 4

depended on Mn and N contents. The effect is relatively small (about

3 FN at 10 wt.% Mn) , however, and not statistically significant at an F

level of 4 (roughly corresponding to a 95% confidence level) . In addition, manganese was often chosen as a ferrite stabilizer when the duplex

stainless steel data (group 10) was evaluated. For these reasons, a Mn

term is not included in the Ni e q.

Since the DeLong and Schaeffler diagrams include coefficients for

Si and Nb , the ACre q

plots for these elements were of particular interest. When welds having reported Nb contents were used (Group 9), the

plot revealed a slope corresponding to a coefficient near 0.7. The

absence of the 0.7-Nb term causes a 2-FN bias, so the term should be

added to the Cre q

to adjust for the Nb effect.

When the ACre q

is examined for a Si effect, a slight trend is also

evident. For equation 4, a coefficient of 0.3 was determined for Si, but

as was the case for Mn, the statistical significance of the Si term was

low. When this coefficient is multiplied by the expected range in Si

content and converted to AFN, the value is near 1 FN. The small statistical significance of this value combined with the limited range in Si

contents caused us to exclude it from the predictive diagram.

The ACre q

plot for Ti was also examined for a coefficient or interaction with N or C. No clear relationships could be resolved.

Although no nonlinear terms could be resolved and used to refine the

Ni and Cr equivalents in the FN equation, the levels at which the scatter

increased (3 wt.% for Mo, 3 wt.% for Mn, and 0.1 wt.% for N) were used to

evaluate the FN diagram.

2 . Mapping

After confirming that they could not be improved by the use of non- linear or interactive terms, the Ni and Cr equivalents were used with a mapping program, which relaxed the restriction of parallel and equispaced

lines. The resulting diagram using the entire database is shown in Fig- ure 3. Straight lines fit well in the Ni e q range of 9 to 17 and in the

Cre q

range of 17 to 30. Above a Ni e q

of 17, there was some evidence of

curvature (dotted lines) , but a scarcity of data prevented deduction of

the true behavior. The appearance is very similar to that of the Schaeffler diagram in Figure 1.

185
18 20 22 24 26 28 30

Creq =Cr+Mo+0.7Nb

Fig. 3. New diagram, produced by mapping program.

The diagram contains equispaced parallel lines between FN values of

4 and 12, explaining the success of the DeLong diagram in the 0 to 18 FN

range with parallel lines. The rapid compression of line spacing above

an FN of 12 and the rotation of the lines above an FN of 35 explains why

the DeLong diagram could not be extended to duplex compositions. The

compression of line spacings and rotation at these higher FN lines ex- plains why the Schaeffler diagram has provided the best estimate of FN in

this region. To determine whether the new diagram represents the FN more accurately than the Schaeffler or DeLong diagrams, quantitative comparisons were made.

B. Evaluation of the New Diagram

The new diagram should have been evaluated using an independent

database. However, all available data were used in the development of

the new diagram. Since the data were gathered from many sources and are assumed to represent current alloy production, they were used to compare

the accuracy of the new diagram to the DeLong and Schaeffler diagrams.

Because these older diagrams have distinctly different ranges, they will

be considered separately.

1. Comparison to the DeLong Diagram (0 to 18 FN)

Since the DeLong diagram was developed for use with the conventional

300- series stainless steel compositions, the measured FN values of the

type 308, 309, and 316 data groups were compared with those predicted by

the DeLong diagram. Relative- frequency histograms of the error distributions are included as Figures 4 through 6. Since each of the groups

186
ERROR INTERVAL

Fig. 4. Relative frequency histogram of the error distribution for the

type 308 data as evaluated by the DeLong diagram.

contains the data for over 100 welds collected from many sources, each is

considered to be representative of current alloy production. The type 308

and 316 data (Figures 4 and 6) exhibit the expected Gaussian error distribution centered on zero with a standard deviation near 2 FN. The type

309 data (Figure 5) exhibit a similar distribution shape and standard

deviation, but they are centered near —2 FN. Such a bias in the mean

indicates that predictions based on the diagram are inaccurate for these

type 309 compositions and do not meet the ± 3 FN accuracy reported for

90% of the data during the development of the DeLong diagram. 8 Separation of the data by source confirmed the same bias for each. Only 43% of

the type 309 data was predicted within a ± 2.5 FN error. For the type

308 and 316 data groups, 82 and 91% of the welds wer-e predicted within a ± 2.5 FN error, respectively. In comparison, the new diagram predicted

88% of the type 308, 91 percent of the type 316, and 92% of the type 3. * data within a ± 2.5 FN error. So, the new diagram is more accurate, es- pecially for the type 309 compositions.

For a broader, more comprehensive evaluation, all the data that had

an FN less than 18 (data group 4) were evaluated. Both the DeLong diagram and the new diagram were used to compute the FN for these compositions. Again, relative- frequency histograms were used to search for : . . or a non-Gaussian shape in the error distribution: The histogram •

187
30

20 - >

O

z

Hi3

O

111 cc

-8 -7 -6 -5 -4 -3 -2 - 1 0 1 2 3 4 5 6 7

ERROR INTERVAL

Fig. 5. Relative frequency histogram of the error distribution

for the type 309 data as evaluated by the DeLong diagram.

were normal; no bias was found. For all the data with an FN less than

18, the DeLong diagram predicted 66% with an error less than 2.5 FN,

whereas the new diagram predicted 84%.

Although the comparisons made between the relative errors found for

DeLong and the new diagram are useful, error magnitudes determined for

the DeLong diagram here should not be compared to- the accuracy statement

(90% of data within ± 3 FN) in that report. 8 The DeLong diagram was developed and evaluated on data from a single laboratory using procedures

more stringent than those in AWS A4.2. This database contains additional

error due to interlaboratory dissimilarities in welding procedures and

specimen preparation techniques. Another- • complication in the evaluation

is that some of the data were outside the Creq boundaries of the published DeLong diagram. Since we chose to evaluate the DeLong diagram

only over its reported range, we excluded these data. Similarly, for

the proposed diagram, computer mapping restrictions limited the Creq

range and excluded a small portion of the data. Manual evaluation of

these data revealed that they fit the extension of the reported lines,

but the scarcity of the data and difficulty of manual calculations caused

us to limit the scope of the diagram to that reported. The number of cases considered for the two diagrams is included in Table 2.

188
40

30 -

>

o

- 8 - 7-6 -5 -4 - 3 - 2-1 0 1 2 3 4 5 6 7

ERROR INTERVAL

Fig. 6. Relative frequency histogram of the error distribution for

the type 316 data as evaluated by the DeLong diagram.

Because this study and others have shown the behavior of some elements to be interactive and therefore the cause of some scatter, the

diagrams were evaluated using data groups 5—8, which had several different combinations of compositional restrictions. Their evaluation is also

summarized in Table 2.

The restriction of the Mo content to less than 3 wt.% resulted in a 2% improvement in the accuracy of the DeLong diagram and a 4% improvement

in the accuracy of the proposed diagram. This significant improvement in

the proposed diagram is the result of either a Cr—Mo interaction, which

is excluded when Mo levels were limited to 3 wt.% during the ACre q analysis, or incomplete dissolution of the high-melting-point Mo contained in

the electrode coating.

Smaller improvements are noted as the N content is restricted to

less than 0.1 wt.% in group 6 or the Mn content is restricted to less

than 3 wt.% in groups 6 and 7. The best fit for both diagrams occurs for

group 8, when all four restrictions (similar to the reported range of applicability of the DeLong diagram) are invoked. However, the improvement: over group 5 is so slight for the proposed diagram that the group 5 re- strictions provide the best combination of accuracy and coefficient simplicity. Scarcity of data at very high element concentration levels do. \

189
suggest some further restrictions: Mn < 10 wt . % and N < 0.2 wt . % . With

these restrictions, the proposed diagram has better than an 88% chance

of predicting the FN with an accuracy of ± 2.5 FN, a value substantially

better than the 68% chance for the DeLong Diagram.

To confirm the validity of these restrictions and to search for

nonlinear behavior, the AFN for the proposed diagram was plotted versus

the various elements. These plots more clearly show the trends seen

during the ANie qand ACre q analyses.

The plot of AFN versus Cr content in Figure 7 has a larger FN range

to include more data and a restricted carbon range to reduce the scatter.

The increased scatter with increasing Mo content is clearly visible.

Either a Cr—Mo interaction or incomplete dissolution of Mo would explain

this scatter. No reduction in scatter was apparent for Mo limits lower

than 3 wt.%. This reinforces the value of restricting the Mo content to

3 wt.% for increased prediction accuracy.

In the plot of AFN versus Mn content in Figure 8, the Mo content

is below 3 wt.%, the FN is below 30, and the Ni content is above 8 wt.%.

The divergence in the Mn response is evident for the two levels of N at

Mn levels above 6 wt.%. At high Mn and N contents, little error in predicted FN values from the new diagram would be expected. For high Mn and

low N contents, however, Mn appears to have a slight ferrite stabilizing

effect. The error is much smaller than that for the Cr—Mo interaction.

Limiting the Mn to 10 wt.% was dictated by the scarcity of data above

this value

.

Table 2. Comparison of the Accuracy of the DeLong Diagram

and the Proposed Diagram for Various Data Groups

.

Data

Group Description

Error Less

than ± 2.5 FN ( %

)

Cases Used

DeLong Proposed DeLong Proposed

4 FN < 18 66 84 747 723

5 FN < 18 and Mo < 3 68 88 706 660

6 FN < 1 8 , Mo < 3 , and

N < 0.1

68 87 643 606

7 FN < 18, Mn < 3 , and

N < 0.1

69 85 635 622

8 FN < 18, Mn < 3,

Mo < 3 , and N < 0 .

1

71 89 601 573

element concentrations in wt.%

190
Fig. 7 Plot of AFN versus Cr content for the new diagram

indicating a Cr—Mo interaction.

TJ

CD

aj

E

ToLU

I z

oCD

Z3

COca

CD

25

20 4

15

10

5

04

-5

-10

-15

-20

-25

FN<30, Mo<3, Ni>8

o N<0.1

N^O.1

° „ 0o °

O o o' fcJJ

.*00 " 00 °v

2 4 6 8 10 12

Manganese Content, wt.%

14

Fig. 8 Plot of AFN versus manganese content for the new diagram

indicating a Mn—N interaction.

191
2. Comparison to the Schaeffler Diagram (0 to 100 FN)

The full range of ferrite content is covered by the Schaeffler diagram (Figure 2) , but this diagram is in volume percent ferrite rather

than FN. So the FN data were converted to volume percent ferrite by the

procedure of Kotecki. 25 Thus we compared the error for the Schaeffler

diagram (using this database converted to percent ferrite) with the error

for the proposed diagram (using this database in FN) . However, comparison of these two relative -frequency histograms is difficult because they

are in different units. To bring the error class intervals into agreement, we assumed that 1% ferrite error is approximately equal to 1.5 FN

error, a reasonable average for the range of the database. In this way, we produced comparable class interval widths to enable a more direct

comparison of error in the two diagrams

.

Since the addition of a N coefficient was one of the factors in

the wide acceptance of the DeLong diagram, we also prepared a modified

Schaeffler diagram containing N. Using DeLong' s coefficient of 30, we fit the Schaeffler diagram to the entire database and obtained a 0.05

wt.% offset term. This meant that the N content should be reduced by

0.05 before being multiplied by 30 and added to the Ni e q.

The proposed diagram and the Schaeffler diagram are compared in

Table 2 for two FN ranges. For FN values greater than 18, the Schaeffler

diagram is substantially less accurate than the proposed diagram. Examination of the error histograms revealed that the error for the new

diagram had a Gaussian shape centered on zero, whereas that for the

Schaeffler diagram was skewed but centered on zero.

For FN values smaller than 18, the Schaeffler diagram is substantially less accurate than the proposed diagram and somewhat less accurate

than the DeLong diagram (Table 3). Again, the error histogram for the

proposed diagram was Gaussian and centered on zero, while the error

histogram for the Schaeffler diagram was centered on zero but skewed.

For FN values smaller than 18, the addition of the N term improved the

accuracy of the (revised) Schaeffler prediction to 70% (greater accuracy

than the DeLong Diagram, but still less accurate than the proposed diagram in this FN range) . Comparison of the three error histograms (proposed, revised Schaeffler with the N term, and DeLong diagrams) indicated

the proposed diagram was characterized by a Gaussian shape centered on

zero, whereas the revised Schaeffler diagram was characterized by a smaller skew to one side, and the DeLong diagram was characterized by a distribution not centered on zero (biased) . It appears that the DeLong

diagram improved on the Schaeffler diagram by adding the N term and

reducing the skewed distribution for FN < 18. The proposed diagram

improves on the DeLong diagram by eliminating the bias and choosing new

coefficients that narrow the standard deviation of the error histogram.

192
Table 3. Comparison of the Accuracy of the Schaeffler Diagram

and the Proposed Diagram.

Data Group

Description

Error

Than ± 9 LessFN (%)

Error Less

Than ± 2.5 FN (%) Cases Used

Schaeffler Proposed Schaeffler Proposed Schaeffler Proposed

FN > 18 35 70 — — 124 125

FN < 18 — — 52 84 771 740

Fig. 9. Solidification-mode data on the coordinate axes of the new diagram.

193
D. Solidification Mode

Hammar 15 and Suutala2

1

evaluated and discussed the use of a predictive equation for the boundaries between the various solidification

modes. Suutala' s solidification data, along with data from Kujanpaa23

and Ogawa 17 were used in this study to evaluate solidification-mode

boundaries for the proposed FN diagram. When the solidification mode was plotted on our Nf e q

and Cre qcoordinate axes, the separation between the

modes was very distinct (Figure 9) , perhaps as good a separation as with

the Hammar -Suutala coordinate axes (Figure 10), which were designed

specifically for solidification-mode separation. This result indicates

that the dependence of high- temperature solidification behavior on com- position is similar to the dependence of the amount of ferrite retained

to room temperature (FN) . Because of the good separation of the solidification modes with our Ni e q

and Cre q expressions, we believe it is premature to conclude that separate Ni e q

and Cre q expressions are necessary

to predict FN and solidification mode in stainless steel welds. Therefore, the boundaries determined for the various solidification zones in

Figure 9 are included on the proposed predictive diagram in Figure 11.

Fig. 10. Solidification-mode data on the coordinate axes

of Hammar and Svenson. 15

194
Fig. 11. New diagram, including solidification-mode boundaries.

The location of the solidification-mode boundaries corresponds well

to changes in the FN response. When solidification boundaries from the

top of the diagram are considered, the solidification boundary of A to AF

(the transition from a fully austenitic structure to a partially ferritic

structure) falls directly on the zero FN line.

The austenite-ferritic to ferritic-austenitic (AF to FA) boundary

denotes the change to a primary ferritic structure, which has been re- lated to a substantial reduction in hot-cracking sensitivity. Using the

fissuring data given by Lundin et al., 1 we noticed that the reported FN

needed to eliminate fissures corresponds very closely to the AF to FA

boundary on the diagram. They reported minimum FN values to prevent fis- sures 1.5 for type 316L (Creq near 20), 2 for type 308 (Cr eq near 18.5),

2.5 for type 316 (Creq near 20.5), 3 for type 308L (Cr eq near 18.8), and

4 for type 309 (Cr eq near 23)

.

This suggests that the AF to FA boundary

could be used as an accurate guide for controlling the hot-cracking tendency, as Kujanpaa also noted in his study. 22

The ferrite—austenite to ferrite (FA to F) transition corresponds to

the beginning of a compression in FN spacing and the start of rotation of

the lines. Microstructurally , this corresponds to a change from a vermicular to a lath structure as the austenitic phase is forced to nucleate

within a ferritic phase.

195
The relationship between FN and solidification mode is shown from

a slightly different perspective in Figure 12. Here, the modified so- lidification equivalents of Suutala 2

1

were used as coordinate axes when

mapping the iso-FN lines. The FN lines are not as straight with these

equivalents, and they do not model compositional effects over as broad

a range as the proposed FN equivalents, but these equivalents indicate

an interaction between FN and solidification mode. We noted a tendency

for the iso-FN lines to move parallel to the boundary between AF and FA

solidification zones and saw indications that iso-FN slopes may differ

in these respective solidification zones.

a

Fig. 12. Schematic representation of iso-FN line shape near

solidification-mode boundaries, indicating a coupling

between the two.

CONCLUSIONS

1. A new predictive diagram for ferrite in stainless steels has been

developed for the' range 0 to 100 FN.

2. The diagram is applicable for Mn contents to 10 wt.%, Mo contents

to 3 wt.%, N contents to 0.2 wt.%, and Si contents to 1 wt.%.

3. The accuracy of the diagram surpasses that of the DeLong and

Schaeffler diagrams in this range for a 950-point database.

4. The diagram includes solidification boundaries that correspond to

changes in the FN response. The AF to FA boundary seems to be a good predictor of cracking sensitivity.

196
ACKNOWLEDGMENTS

Several of the authors received financial support from the Department of Energy and the Welding Research Council. The authors are also

indebted to Alloy Rods Company, Teledyne McKay, Sandvik Steel Company,

Welding and Wire Division, Smitweld, Soudometal, and the Welding Institute, who supplied data.

REFERENCES

1. Lundin, C. D. , DeLong, W. T. , and Spond, D. F. 1975. Ferrite—

Fissuring Relationship in Austenitic Stainless Steel Weld Metals.

Welding Journal 54. pp . 241s—246s.

2. Baeslack, W. A., Ill, Duquette, D. J., and Savage, W. F. 1978.

Technical Note: Stress Corrosion Cracking in Duplex Stainless Steel

Weldments. Welding Journal 57. pp . 175s—177s.

3. Siewert, T. A. 1978. How to Predict Impact Energy from Stainless

Steel Composition. Welding Design and Fabrication 51. pp . 88—90.

4. Szumachowski , E. R. and Reid, H. F. 1978. Cryogenic Toughness of

SMA Austenitic Stainless Steel Weld Metals: Part 1 — Role of Ferrite. Welding Journal 57. pp. 325s—333s.

5. Read, D. T. , McHenry, H. I., Steinmeyer, P. A., and Thomas, R. D.,

Jr. 1980. Metallurgical Factors Affecting the Toughness of 316L

SMA Weldments at Cryogenic Temperatures. Welding Journal 59.

pp. 104s—113s.

6. Siewert, T. A. 1986. Predicting the Toughness of SMA Austenitic

Stainless Steel Welds at 77 K, Welding Journal 65. pp . 23—28.

7. Kotecki, D. J. 1987. Ferrite Control in Duplex Stainless Steel

Weld Metal. Welding Journal 66. pp. 273s—278s.

8. DeLong, W. T. 1974. Ferrite in Austenitic Stainless Steel Weld

Metal. Welding Journal 53. pp. 273s-286s.

9. Schaeffler, A. 1949. Constitution Diagram for Stainless Steel Weld

Metal. Metal Progress 56. pp. 680—680B.

10. Olson, D. L. 1985. Prediction of Austenitic Weld Metal Microstructure and Properties. Welding Journal 64. pp . 281s—295s.

11. Hull, F. C. 1973. Delta Ferrite and Martensite Formation in Stainless Steels. Welding Journal 52. pp . 193s—203s.

12. Espy, R. H. 1982. Weldability of Nitrogen-Strengthened Stainless

Steels. Welding Journal 61. pp. 149s—156s.

197
13. Szumachowski , E. R. and Kotecki , D. J. 1984. Effect of Manganese

on Stainless Steel Weld Metal Ferrite. Welding Journal 63.

pp . 156s—161s.

14. McCowan, C. N. , Siewert, T. A., Reed, R. P., and Lake, F. B. 1987.

Manganese and Nitrogen in Stainless Steel SMA Welds for Cryogenic

Service. Welding Journal 66. pp . 84s—92s.

15. Hammar, 0. and Svenson, U. 1979. Solidification and Casting of

Metals . The Metals Society. London, pp . 401—410.

16. Mel'Kumor, N. and Topilin, V. V. 1969. Alloying Austenitic Stainless Steel with Nitrogen. Obrabotka Metallor Aug. pp . 47—51.

17. Ogawa, T. and Koseki, T. 1988. Weldability of Newly Developed

Austenitic Alloys for Cryogenic Service: Part II — High-Nitrogen

Stainless Steel Weld Metal. Welding Journa l 67. pp . 8s—17s.

18. Okagawa, R. K. , Dixon, R. D., and Olson, D. L. 1983. The Influence

of Nitrogen from Welding on Stainless Steel Weld Metal Microstructure. Welding Journal 62. pp. 204s—209s.

19. Kotecki, D. J. 1986. Silicon Effect on Stainless Weld Metal Ferrite. IIW. Doc. II-C-779-86. The American Council of the

International Institute of Welding. Miami.

20. Takemoto, T. , Murata, Y. , and Tanaka, T. 1987. Effect of Manganese

on Phase Stability of Cr—Ni Nonmagnetic Stainless Steel. In High

Manganese Austenitic Steels . Ed R.A. Lula, ASM Conference Proceedings, Oct. 10—15, 1987, ASM International, Cincinnati, Ohio.

pp. 23-31.

21. Suutala, N. 1982. Effect of Manganese and Nitrogen on the Solidification Mode in Austenitic Stainless Steel Welds. Metallurgical

Transactions A 13(12). pp. 2121—2130.

22. Lippold, J. C. and Savage, W. F. 1982. Solidification of Austenitic Stainless Steel Weldments: Part III — The Effect of Solidification Behavior on Hot Cracking Susceptibility. Welding Journal

61. pp. 388s—396s.

23. Kujanpaa, V. 1985. Role of Steel Type and Impurities in Solidification Cracking in Austenitic Stainless Steel Welds. Metal Construction 17(1) pp. 40R—46R.

24 . Standard Procedures for Calibrating Magnetic Instruments to Measure

the Delta Ferrite Content of Austenitic Stainless Steel Weld Metal . 1986. ANSI/AWS A 4.2-86, American Welding Society. Miami.

25. Kotecki, D. J. 1982. Extension of the WRC Ferrite Number System.

Welding Journal 61. pp . 352s—361s.

198
26. McCowan, C. N. , Siewert, T. A., and Olson, D. L. 1988. Stainless

Steel Welds: Prediction of Ferrite Content, Interagency report,

National Bureau of Standards, Boulder, Colorado; in preparation.

27. David, S. A., Vitek, J. M. , and Hebble, T. L. 1987. Effect of

Rapid Solidification on Stainless Steel Weld Metal Microstructures

and Its Implications on the Schaeffler Diagram. Welding Journal 66.

pp . 289s—300s.

199

MECHANICAL PROPERTIES OF ELECTRON BEAM

WELDS IN THICK COPPER

%

T. A. Siewert and D. P. Vigliotti

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

Electron beam welding was used to make similar and dissimilar butt joints in 25-mm-thick plates of copper alloys

C10700 and C17510. The mechanical properties of these joints

were measured at 76 and 298 K using reduced- section transverse

tensile specimens. Elongation measurements of a 25 -mm- long

gage length consisting of weld, heat-affected zone (HAZ) , and

base metal were misleading. Therefore, for some specimens,

elongation was measured on 1-mm lengths, a distance more characteristic of the weld and HAZ widths. These measurements

revealed that the first strain occurred in the soft weld and

HAZ. The effect of flow- localization constraint on the soft

regions was small for the specimens tested, which ranged from

15 to 128 mm2 in cross-sectional area.

INTRODUCTION

The design for the Compact Ignition Tokamak (CIT) could require

welds in 25-mm-thick sheets of copper alloys. Copper alloy C10700

(Cu—0.08 mass % Ag) , one candidate material, has a good combination of

strength and conductivity, especially when work hardened. In regions

of the structure where stresses could exceed the yield strength of alloy

C10700, precipitation-hardenable copper alloy C17510 (Cu—0.4 mass % Be—

1.8 mass % Ni) is being considered. To join the copper conductor segments into a continuous conductor, welds might be necessary for the

following alloy pairs: C10700 to C10700, C10700 to C17510, and C17510

to C17510

.

Copper alloys have been joined successfully for many years by the

common welding processes, including shielded metal arc, gas metal arc,

and gas tungsten arc welding. 1 Such welds are produced with high heat

inputs and high preheat temperatures to overcome the effects of the h i ;-.h

201
Table 1. Material Characteristics and Properties*

Alloy C10700 Alloy C17510

Heat 9049 50187

Yield strength (MPa) 304t 794

Tensile strength (MPa) 321 845

Elongation (%) 22 12

Hardness

Composition (mass %)

96 Brinnell 102 Rb

Be —0.38

Ni -

1.79

Ag approx. 0.09 —

Thermomechanical

history

cold work: 40% cold work: 37%

,

followed by

age hardening

at 482°C for 2 h

reported by the supplier

t determined by the 0.5%-strain (elastic and plastic) offset technique;

estimated to be similar to data measured by the 0.2%-strain (plastic

only) offset technique

Table 2. Welding Conditions for EB Welds

Voltage (kV)

Current (mA)

Power (kW at workpiece)

Travel speed (mm/min)

Focal plane (mm from top of plate)

Orientation

Oscillation (mm)

Pass Sequence

55

350

19.3

380

20

beam horizontal,

normal to plate

1 single pass, one side

202
thermal conductivity of the copper alloys. This high heat input con- tributes to a loss in strength in the heat- affected zone (HAZ) through

recrystallization of the cold-worked structure in alloy C10700 or solution heat treatment of the precipitation-hardened structure in alloy

C17510

.

2

This study evaluated a high-energy beam welding process, electron

beam (EB) welding, because the energy density enables it to produce a weld with a significantly narrower bead and HAZ width without preheating.

These attributes are important: the elimination of preheating avoids the

cost and safety considerations involved in bringing a joint to 500°C, and

the narrower bead and HAZ contribute to a higher effective weld strength

through flow localization. 3

Another high-energy beam welding process, laser welding, was to be

compared to the EB process. Unfortunately, acceptable welds could not be

produced within the time constraints of this project. Although the C0 2 laser has a power of 12 kW, its wavelength was reflected strongly by the

copper, so the beam could not penetrate the joint. Various attempts were

made to reduce the section by beveling and later filling the bevel with

molten copper, but these welds had an unacceptable level of porosity.

In addition to electrical conductivity and strength requirements,

the welds must withstand thermal cycles. The structure is cooled initially by liquid nitrogen, but electrical resistance warms the structure

to room temperature. To provide data for the design of the structure,

the mechanical properties of these material combinations were evaluated

at the extremes of this temperature range.

MATERIALS AND METHODS

Alloys C10700 and C17510 were from heats that met the compositional

requirements of the Copper Development Association and the mechanical

property requirements for CIT applications. Detailed information on the

material is listed in table 1. Both materials were received in the form

of 25-mm- thick plate.

Since the amount of plate was limited, reduced- section specimens

were chosen for the tensile . tests . The plates were sawn into 50-mm x 150-mm blocks, then welded along the 150-mm edge to form 100-mm x 150-mm

x 25-mm weldments. The welding parameters are listed’ in table 2.

Tensile specimens were removed from the EB weldments and pulled

to failure in uniaxial tension to evaluate the effect of welding on the

mechanical properties. The tests were performed in a modified load frame

that had a reaction cylinder which enabled placing a Dewar of liquid nitrogen around the specimen. The test procedure and the design of this

load frame have been reported elsewhere. 4 The test matrix included two

temperatures, three material combinations at the weld, and a variety of

cross-sectional areas.

203
Normally, properties of welds are measured on tensile specimens

oriented along the length of the weld and consisting only of weld metal.

The narrowness of the EB welds precluded this specimen orientation and

required transverse specimens. However, the transverse orientation is more desirable, since it will be the primary loading direction in the

coil. In this orientation, the gage lengths of the standard round- and

flat- section specimens were not homogeneous but included a variety of

materials: starting from one gage mark, the gage length consisted of

base metal, HAZ, weld metal, HAZ, and the second base metal. Because

each of these regions has different properties, the mechanical properties must be interpreted on the basis of the specimen's composite nature.

Two of the most important effects of the composite nature are flowlocalization constraint and localized yielding.

Constraint

The HAZ regions, weakened by recovery and recrystallization of the

grains during welding, affect the joint strength. These regions, each

containing approximately 5 mm of the tensile specimen's 50 -mm gage

length, should strongly affect the tensile properties of a very thin

( 1 -mm-diameter) tensile specimen. Conversely, because of the constraint

of the adjacent base metal (Poisson's ratio effect), this HAZ should have

less effect on a thicker, say 25 -mm- diameter , tensile specimen. To mea- sure this effect, a series of specimen cross sections was tested. With

the limited material available, we were able to produce cross-sectional

areas ranging from 15 to 128 mm2

. The 128-mm2 cross section was that of

a round specimen with a diameter of 12.8 mm. This diameter was expected

to cause only marginal constraint in the 11 -mm- thick weldment (5 -mm HAZ,

1-mm weld, 5-mm HAZ-softened region), but it was the largest diameter

obtainable from the small weldments.

Localized Yielding

Because the weld and HAZ have less strength than the base metal, the

initial strains occur in these weaker regions until they reach the yield

strength of the base material. Strain measurements at a series of marks

1 mm apart were used to determine the strain history for the various

regions of the gage lengths in several specimens.

A bench- top lathe was used to scribe the marks on the tensile sur- face after it was coated with machinist's layout ink. When the specimen

was placed in the tensile fixture, the marks were visible in the eyepiece

of a traveling microscope, which had been mounted on the frame of the

tensile machine.

By monitoring the stress—strain curve, we could stop the test at

selected strain levels for measurement of the relative movement of these

1 -ram- spaced marks, which enabled determination of the local strain in the

base metal, HAZ, and weld metal.

204
RESULTS AND DISCUSSION

Microstructure

The welds were initially evaluated by removing transverse sections

for micros tructural evaluation. Figure 1 shows the etched cross section

of the EB weld of the C10700—C10700 material combination. The weld

fusion zone was approximately 1 mm wide and uniform over the 25 -mm thick- ness of the joint. It was abutted on each side by a coarse-grained HAZ

about 1 mm thick, then a fine-grained HAZ about 4 mm thick. Beyond the

fine-grained HAZ was the base metal. Figure 2 shows the etched cross section of the EB weld joining alloy C17510 to C17510. The weld fusion

zone was approximately 1-mm wide and uniform over the 25 -mm thickness of

the joint. The weld was abutted on each side by a visible HAZ, 6 mm in

Figure 1. Cross section of alloy C10700-alloy C10700 weld;

nitric and lactic acid etch.

205
10mm 20

Figure 2. Cross section of the alloy C17510—alloy C17510 weld;

nitric and lactic acid etch.

width. Figure 3 shows the etched cross section of the EB weld joining

alloy C10700 to C17510. The different interaction of the two materials

with the electron beam caused a deflection of the beam within the joint.

The early, procedural development welds showed the beam centered in the

joint at the surface, but deflected about 1 or 2 mm into alloy C10700 at

the bottom. This was corrected by angling the beam toward alloy Cl 75 10.

For this reason, the weld interface in figure 3 shows some curvature.

Its width is similar to those of previous welds. The HAZ width is

different on each side of the joint. The HAZ on the right matches the

microstructure of the HAZ in figure 1; the haz on the left, that in

figure 2.

206
Illljllll lllllllll llllllll

10mm 20

Figure 3. Cross section of the alloy C10700—alloy C17510 weld;

nitric and lactic acid etch.

Hardness Tests

The polished cross section shown in figure 1 was evaluated for

hardness by producing three hardness traverses across the weld at the

top, midthickness, and bottom of the weld. The hardness was measured

with the Rockwell F technique at 3 -mm intervals; results are included in

table 3. The hardness data on the base metal agree well with the data

provided by the suppliers (table 1). On the basis of the general relationship between hardness and strength, the substantial reduction in

hardness in the HAZ and weld should indicate a similar reduction of

strength in these regions.

Tensile Tests

The tensile test results are summarized in table 4. The yield

strength was determined by the 0 . 2%-strain-offset method, a method that

assumes uniform elongation in the 25-mm gage length. The 0.5% -strain

value measured under load (combined plastic and elastic strain) has traditionally been used for copper alloys. We chose the 0.2% -strain value

207
Table 3. Hardness Traverses across EB C10700—C10700 Weld

Location

(mm)

Description

Hardness (Rp)

Top Middle Bottom

-9 Base metal 83 84 83

-6 HAZ—BM interface 56 49 48

-3 Fine HAZ 38 39 37

0 Weld 16 32 26

+3 Fine HAZ 36 37 37

+6 HAZ—BM interface 57 49 48

+9 Base metal 84 82 83

+12 Base metal 82 84 84

Table 4. Mechanical Property Data

SpecimenMaterial

Couple

T

(K)

Crosssectional

Area

(mm2

)

Yield

Strength*

(MPa)

Tensile

Strength

(MPa)

Elongationf

(%)

Reduction

of Area

(%)

EB - 1 Cu—Cu 298 32 89 222 22 68

EB-2 Cu—Cu 76 32 187 376 23 68

EB - 3 Cu—Cu 298 15.5 74 181 27 50

EB-4 Cu—Cu 76 15 132 319 36 75

EB - 5 Cu—CuBe 298 15.5 87± 158* 17* 46*

EB - 6 Cu—CuBe 76 15.5 144? 261* 14* 42*

EB - 7 Cu—Cu 298 128 59 206 49 75

EB - 8 Cu—Cu 76 128 70 328 41 69

EB-9 CuBe—CuBe 298 126 253 457 16 58

EB- 10 CuBe—CuBe 298 125 251 454 15 ' 60

EB - 1

1

CuBe—CuBe 76 126 320 593 13 53

EB - 12 CuBe—CuBe 76 126 318 591 14 47

EB-13 Cu—CuBe 298 126 74 225 45 90

EB- 14 Cu—CuBe 298 126 87 232 34 74

EB - 15 Cu—CuBe 76 126 98 359 36 77

EB-16 Cu—CuBe 76 126 105 356 46 89

107-1 Cu 298 126 280 308 40 86

107-2 Cu 76 126 356 427 56 82

175-1 CuBe 298 126 720 805 17 19

175-2 CuBe 76 126 795 — — —

based on 0 . 2% - strain-offset method

t based on 25 -mm gage length

^ lower than expected, from preliminary weld with flaws

(see EB-13 through EB-16 for second weld)

208
(plastic strain only), which is traditionally used for steel, since it

is now favored. The transition from elastic to plastic strain is quite

gradual for these specimens (see figure 4), and the two techniques (0.2%

plastic and 0.5% total strain) provide similar estimates of the loadcarrying capacity for this material.

Since this strain was not uniformly distributed in the weld specimens, the actual strain in the weld was greater. Figure 5 shows the

development of strain at 1-mm intervals down the length of specimen EB-3.

The first measurement was at an average strain of 10%. Figure 5 also indicates the locations of strain values in the weld and HAZ regions. The

localization of the initial strain at the weld and two HAZs is apparent.

As this material strains, it work-hardens and increases in strength.

When it reaches the strength of the base metal, it also begins to strain

plastically.

The amount of strain that occurs in the weld and HAZs is expected to

be a function of the specimen's cross-sectional area. Increasing cross- sectional area results in greater constraint: the stronger base metal

prevents plastic flow and work hardening in the softer weld and HAZs.

Such behavior can be examined in a more quantitative manner by plotting

the 0.2% yield strength as a function of cross-sectional area. However,

figures 6 and 7 do not show any effect of constraint on yield strength

for the cross-sectional areas included in this study.

Limited data linking hardness to strength for copper alloys are available. 5 The identities of these alloys are not available, but the

data are probably a collection of various alloys, grain sizes, and

amounts of cold work. These data, shown in figure 8, can be used to

relate the microstructure and hardness to the mechanical properties of

the alloy C10700 welds. The 298-K yield- strength data for the transverse

specimens can be correlated to the 80-MPa yield strength predicted by

the Rg-38 hardness of the fine-grained HAZ. The very soft (Rg-16) weld

region should have caused the initial yielding to occur near 11 MPa.

Therefore, the constraint of the specimen geometry was sufficient to

prevent the 1-mm-wide weld from deforming, and the deformation occurred

in the fine-grained HAZ. If the specimen cross section is increased

further, the fine-grained HAZ will be constrained, and the yield strength

should approach the base-metal value of 280 MPa.

The fact that the deformation was localized near the weld causes

us to reconsider the applicability of a 0.2%-strain design limit. The

0.2%-strain limit prevents massive deformation in a large volume. However, for these welds, the plastic strain was limited to a region on the

scale of millimeters, until the base metal yield strength was reached.

Perhaps a 0.5 or 1% plastic strain in this region can be tolerated,

especially for materials, such as these copper alloys, that have good

work-hardening characteristics (figure 4). Table 5 lists the specimen

strengths for differing plastic strains over the 25-mm gage length. The

yield strength of the specimens increased by an average of 30% as the

plastic strain was increased from 0.2 to 2%.

209
Figure 4. Representative stress-strain curve for alloy C10700—alloy C10700

weld showing substantial ductility and strain hardening.

Figure 5. Elongation at 1-nun intervals during the straining of

specimen EB-3 (measured at two levels of strain).

210
Strength

(MPa)

Strength

(MPa)

400

300

200 -

100

• - yield (0.2%)

X - tensile 298 K

Cu alloy 10700

“Nrh

20 40 60 80 100 120 140

Cross-sectional Area (mm2

)

Nr+J

Base

Metal

Figure 6. Strength versus cross-sectional area for

alloy C10700—alloy C10700 weld at 298 K.

400 -

300 -

200 -

100 -

• - yield (0.2%)

X - tensile 76 K

Cu alloy 10700

20 40 60 80 100 120

Cross-sectional Area (mm2

)

Nr

140

NH-I

Base

Metal

Figure 7. Strength versus cross-sectional area for

alloy C10700—alloy C10700 weld at 76 K.

211
Figure 8. Yield strength (0.5% extension under load) versus

hardness for a variety of copper alloys at 298 K. s

Table 5. Yield Strength as a Function of Strain

Tempera- Yield Strength (MPa)

Specimen ture Area

(K) (mm2

) 0 . 2%* 0 . 5%* 1%* 2%*

EB-1 298 32 89 100 109 125

EB-2 76 32 187 198 205 231

EB-3 298 15.5 74 80 88 103

EB-4 76 15 132 150 166 196

EB-7 298 128 59 70 83 101

EB-8 76 128 70 85 99 122

* plastic strain only

212
Both the weld and the base metals became stronger at reduced temperatures. Figure 9 portrays the data for the alloy C10700—alloy C10700

welds and base metal. The weld data are the average of the three tests

at each temperature. Upon cooling from 298 to 76 K, the weld strength

increased approximately 50%, and the base-metal strength increased

approximately 25%.

Figure 9. Yield strength versus temperature for alloy C10700—

alloy C10700 welds and 40% cold-worked base metal.

CONCLUSIONS

1. The weld yield strengths were less than half those of the base

metals

.

2. Much of the early plastic strain during tensile testing is re- stricted to the weld and HAZ because the heat of welding softens

the cold-worked structure.

3. Owing to the wide, fine-grained HAZs in the welds between alloys

C10700 and C10700, few constraint effects were seen for the specimen

sizes included in this study.

213
ACKNOWLEDGMENTS

This program was funded primarily by the Office of Fusion Energy,

Department of Energy. The production of the welds was funded by Princeton Plasma Physics Laboratory. The authors acknowledge the assistance of

C. N. McCowan in the tensile testing of the specimens and of D. Shepherd

in the metallography and hardness testing. The EB weld was produced by

Sciaky Brothers.* Lawrence Livermore National Laboratory supplied the

C17510 copper alloy; Princeton Plasma Physics Laboratory, the C10700

copper alloy.

REFERENCES

1. Welding. Brazing, and Soldering . vol. 6 of the Ninth Edition of the

Metals Handbook . American Society for Metals, Metals Park, Ohio

(1983), p. 400.

2. Ibid , p. 418.

3. Mechanical Testing , vol. 8 of the Ninth Edition of the Metals

Handbook . American Society for Metals, Metals Park, Ohio (1985),

p. 573.

4. D. T. Read and R. L. Tobler, "Mechanical Property Measurements at

Low Temperatures," in Advances in Cryogenic Engineering — Materials , vol. 28, Plenum Press, New York (1982), pp. 250-268.

5. C. Bushnell, Princeton Plasma Physics Laboratory, Princeton, New

Jersey, private communication (1987).

This company is listed for information only. No endorsement by NBS is

implied.

214
CRYOGENIC MATERIAL PROPERTIES

OF STAINLESS STEEL TUBE -TO -FLANGE WELDS*

T. A. Siewert, C. N. McCowan, and D. P. Vigliotti

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The mechanical properties of stainless steel beam tube -to

flange welds for a cryogenic piping application were measured.

A planar specimen was developed to duplicate the constraint,

loading, and heat-sink properties of the circular joint while

reducing preparation time and cost. We evaluated specimens

containing welds between the tube material (21Cr—6Ni—9Mn)

and the three stainless steels being considered for the flange

materials: AISI types 304L and 316L and 21Cr—6Ni—9Mn. The me- chanical property tests consisted of three phases: simple tensile testing to failure, tensile testing of notched specimens

(where the notch simulated fabrication flaws) , and fatigue

testing of notched specimens for the 4 x 10 4 cycle design life

of the structure. Flanges of type 316L stainless steel produced welds with the best combination of strength and ductility

in all three phases of testing and at 295 and 4 K.

INTRODUCTION

Their very high current-carrying capacity has made 4-K superconductors the preferred choice for applications requiring powerful magnetic

fields. Thin-walled stainless steel tubes at the center of the superconducting magnets will contain the proton and antiproton beams. These

"beam tubes" are connected to the system by welds of their flanges. A

failure in the flange welds would disable the system. This study provides the information needed for design of reliable beam tube- to-flange

welds

.

Various stainless steel alloys have been characterized for cryogenic

service. 1 When these results were reviewed, austenitic stainless steel

alloys had some of the best combinations of strength and toughness at

4 K. 2

* Submitted to Journal of Engineering Materials and Technology

.

215
An understanding of the role of nitrogen in determining the strength

and toughness is necessary before appropriate combinations can be chosen

for tube- to -flange welds. Nitrogen additions substantially increase the

cryogenic strength of stainless steels, but at the same time, they re- duce the fracture toughness. This reciprocal relationship is shown in

Figure 1 for types 316L and 316LN stainless steels containing different

amounts of nitrogen (N indicates intentionally added nitrogen). 2

Figure 1 also includes data for welds of matching composition. 3 Their lower strength—toughness combinations indicate that, for equal

weld-and base-metal strengths, welds are the fracture-critical portion of

the structure.

These strength—toughness data for type 316 stainless steel welds

were developed for thick, multipass welds because they will be used in

joining the segments of large magnet structures, and they have dimensions

suitable for the standard 25 -mm- thick compact- tension fracture - toughness

specimen. A literature search revealed no cryogenic data for the singlepass welds used to join stainless steel tubes to flanges. Since the

properties for single-pass welds could differ substantially from those

for multipass welds, we determined the 4-K strength and ductility of

single-pass welds of the three flange materials under consideration.

YIELD STRENGTH, ksi

c

05

*

CO* CO

Hi

z

X

o

D

o

IDGCD

H

O

<

GC

LL

Figure 1. Yield strength versus fracture toughness for 316 type alloys

at 4 K. Scatter bands represent one standard deviation.

216
MATERIALS

Three types of stainless steels were evaluated for use as the flange

material. All were purchased from commercial sources with certificates

of their room- temperature mechanical properties and composition, which

are included in Table 1.

The alloy identified as 21Cr—6Ni—9Mn meets the cryogenic strength

requirements for the application. It is included in the list of potential flange materials to provide a reference for the matching-composition

joint. Although Figure 1 reports data for only types 316L and 316LN

base-metal compositions, the inverse relationship between strength and

toughness seems to apply to other austenitic alloys as well. Therefore,

since this 21Cr—6Ni—9Mn alloy has a much higher nitrogen content than the

other alloys in Table 1, its toughness is expected to be lower. Since

measurements of fracture toughness are difficult for this joint configuration and thickness, we have attempted to compare the materials in

terms of relative ductility, as measured by plastic strain to failure.

Type 304L stainless steel is included because it is probably the

most common austenitic stainless steel. Its low cost, ready availability, and ample cryogenic mechanical property data (for multipass

Table 1. Mechanical Properties and Compositions of the Three Types

of Stainless Steel at Room Temperature

Material

Specification

Designation

21Cr—6Ni—9Mn

ASME SA-412

Type XM-11

304L

ASME SA- 240

Type 304L

316L

ASME SA- 240

Type 316L

316LN

ASME SA- 276

Type 316LN*

Thickness (mm) 9.5 9.5 9.5

Heat number 65724 Z1136 C2713 —

Yield strength (MPa) 427 322 357 —

Tensile strength (MPa) 749 588 609 —

Elongation (%) 50 56 52 —

Reduction of area (%) 62 63 69 —

Hardness (Brinell) 205 150 176 —

Composition (wt.%)

C 0.023 0.018 0.014 0.02

Mn 9.16 1.69 1.50 1.50

Si 0.35 0.53 0.54 0.50

P 0.030 0.029 0.015 0.03

S 0.001 0.007 0.005 0.02

Cr 19.76 18.30 16.60 17.0

Ni 7.14 9.27 10.88 11.0

Mo — — 2.10 2.5

N 0.32 0.079 0.084 0.13

Nominal values

217
welds) 4 make it a reference material to which the other steels can be

compared. It has lower strength and higher toughness than 21Cr—6Ni—9Mn

.

Type 316L stainless steel is included because it has one of the best

strength—toughness combinations. 2 ’ 4 A type 316LN alloy was used in the

first phase of testing because it was available from a previous program.

Difficulty in obtaining more of this alloy for the second phase of testing forced the switch to type 316L. The primary difference in these

alloys is nitrogen content—0.08 wt.% in type 316L and 0.13 wt.% in

type 316LN.

The effect of substituting 316L for 316LN is expected to be small

for this application. The nitrogen content in an autogenous (without

filler metal) gas tungsten arc (GTA) weld is approximately the average

of the nitrogen contents of the two joint materials, since each contributes about half the metal that fuses to form the weld bead. Thus, the

expected nitrogen content of the joint between the 316L flange material

and the 21Cr—6Ni—9Mn tube is 0.20 wt.%, whereas that for the joint between the 316LN flange material and the 21Cr—6Ni—9Mn tube is 0.22 wt.%.

The 0.02-wt.% difference is considered to be insignificant because the

expected nitrogen content of the joint of 21Cr—6Ni—9Mn flange material

and tube of like composition is 0.32 wt.%.

These three flange materials were joined to the 2lCr—6Ni—9Mn tube

with a 0.9 -mm thick wall. Although the mechanical properties of 21Cr—

6Ni—9Mn have been extensively characterized at room temperature, little

information was available on its mechanical properties at 4 K, especially

for tubing formed by either extrusion or rolling and seam-welding pro- cesses. During this test program, lack of this information plus the introduction of a new tube with different wall thickness prompted a cursory

investigation of these properties of the 21Cr—6Ni—9Mn tube, which is

reported in Appendix A. Although the difference in wall thickness and

type of weld should not significantly alter our test results, the test

program was completed without the introduction of these variables.

Figure 2. Cross section of a typical tube -to -flange joint.
PROCEDURE

Specimen Design

Tube- to -flange welds for cryogenic piping tend to be single-pass

welds that are performed autogenously with gas tungsten arc welding.

The tube is typically seated inside the flange where a circumferential

weld forms a joint that provides mechanical strength and an air-tight

seal. The schematic for the joint between a tube with a 0.9 -mm- thick

wall and a much heavier flange section is included in Figure 2. However, this type of joint is expensive to produce in the quantities

needed for this test program, and reducing the specimen cost by slicing

the joint into strips results in a bending stress when the joint is

loaded.

To retain the differences in section mass while avoiding the bending moment, the test specimen configuration shown in Figure 3 was developed. Specimens of this design were produced rapidly and inexpensively while duplicating the critical characteristics of the tubeto-flange weld; they were suitable for standard methods of mechanical

testing. 5 To simulate the fact that most of a circular tube- to -flange

weld is free from starts or stops, wider specimens were produced, and

the weld start and stop regions were removed. The welding was performed at 100 A with a travel speed of 12 cm/min. Figure 4 shows the

welding fixture with clamps to maintain the alignment and with the gun

held in a motorized carriage. When several lack-of- fusion flaws were

found in the phase 1 specimens, the electrode tip angle was decreased

from 60° to 30°; the problem did not occur in subsequent specimens.

Figure 3. Schematic diagram of the specimen showing the

four notch locations: 1 - flange HAZ; 2 - weld;

3 — tube HAZ; 4 — tube.

219
Figure 4. Welding of the specimens in an automated fixture.

Test Plan

To evaluate the sensitivity of the joints to the expected failure

conditions, they were subjected to three test phases:

1. Uniaxial tensile testing to measure the joint strength

2. Uniaxial tensile testing of specimens with notches to simulate the

effect of fabrication flaws and the inherent notch in the joint

design on strength and ductility

3 . Fatigue loading of notched specimens to evaluate the fatigue

resistance to growth of fabrication flaws

Since each subsequent test phase increased in complexity and cost, corre- spondingly fewer material combinations were used.

The tensile- loading evaluation, phase 1, included all three flange

materials joined to the 21Cr—6Ni—9Mn tube. Figure 5 shows a specimen in

a room- temperature load frame. The two materials with the best strength

and elongation results advanced to phase 2, where each was notched in

four locations: the flange heat-affected zone (HAZ) , the weld, the tube

220
HAZ, and the tube, as illustrated in Figure 3. The phase 2 specimen con- figuration was the same as phase 1, except a 60° included- angle notch was formed on each edge of the specimen at the desired locations. Each side was notched to a depth of 10% of the specimen width and a root radius of

0.064 mm. Figure 6 shows a typical notch, as photographed on our optical

comparator. For this particular notch, the comparator calculated a root

radius of 0.066 mm.

On the basis of the notch- tensile test results, the four most promising material-notch combinations advanced to the fatigue testing of

phase 3. In phase 3, we determined the fatigue life of the tube, tube

HAZs , welds to type 304L flanges, and welds to type 316L flanges. The

specimen configuration was identical to that in phase 2. Specimens were

tested in sets of five to enable a meaningful S—N curve to be drawn.

Because both the strength and ductility of the stainless steels are affected substantially by temperature, the joints were evaluated at two

temperature extremes, 4 and 295 K. The 4-K tests were performed with the

specimens immersed in liquid helium. Details of the 4-K load frame have

been reported elsewhere. 6

SP^‘

~ '

Figure 5. Specimen in a room- temperature load frame.

221
tVl'irG—: • ~ ’“V* ^{W'c^Vt^yV -

r M .*- £02&

QBfLlSS I

-• yt w Off

..• i: '.

Vi V -— fi i-j t-j

, «» 1 ; t*’l C

: j i Z • it PiK]

'. * ~ a *

Figure 6. Examination of the notch radius on an optical comparator.

Material Property Measurements

For this study, the yield strength was defined as the load divided

by the original cross-sectional area, where the value of the load is that

when the testing machine crosshead had been displaced 0.05 mm beyond

simple elastic loading of the specimen. The displacement of 0.05 mm is

equivalent to an average of 0.2% plastic strain over the nominal 25 -mm specimen gage length (distance from flange to reinforcing welds on the

spacers)

.

The term fracture strength is used when failure occurs as the

stress—strain curve is increasing; the term tensile strength, when fail- ure occurs as the stress—strain curve is decreasing or when it has asymptotically reached a constant value. The term 0.2% yield strength is not

accurate in this case because the composite nature of the gage length

(weld + HAZ + base metal) encourages nonuniform, localized yielding. 7

222
Therefore, in this study, fracture strength was used because the strength

was influenced more by the geometry of the specimen than the inherent

material response; it was defined as the maximum load divided by the

original cross-sectional area after notching.

A substantial range in ductility was observed among the specimens.

The composite nature of the gage length, however, precluded reporting

this ductility in units of uniform strain. If the elongation had been

uniform, this displacement could have been converted to strain by dividing by the gage length. Instead, the crosshead displacement of the testing machine was used. The displacement measurements were corrected to

purely plastic extension. The elastic component was removed by measuring

the displacement from a line parallel to the initial elastic loading

slope. Plastic extension measurements provide an excellent measure of

the relative ability of the joints to tolerate strain.

For fatigue testing, the pin- loaded specimen design required a minimum tensile load to eliminate vibration in the load train. Even for the

lowest load levels, this minimum was never greater than 10% of the maximum load for any of the tests. The minimum load, rather than AP (the

difference between the maximum and minimum loads) , was kept constant

throughout the tests to obtain the largest AP value for each test. This

choice produces a more conservative value of the fatigue life. We did

not evaluate the effect of multiple cracks in the specimens (four notches

per specimen, as shown in Figure 3)

.

Since the structure is designed for 4 x 10 4 load cycles, stress

at this life must be estimated to evaluate the various materials. We

determined the life for several specimens near the desired range. Then

we tested several more at a higher load, which resulted in shorter life)

.

By drawing a smooth line through the average life at these load levels, we were able to determine the stress at which this line intersected

4 x 10 4 cycles and, thus, to evaluate the various materials and notch

locations

.

RESULTS AND DISCUSSION

Tensile Tests (Phase 1)

Uniaxial tensile tests were conducted on welds of the three flange

materials at both 295 and 4 K.

A. Tensile Tests at 295 K

The yield strengths of the three types of specimens are very similar, as shown in Table 2, indicating that differences in nitrogen con- tents of the flange materials had little influence on the yield strength:;

of the joints at this temperature. Strengthening due to constraint effects in the welds exceeded the small differences in nitrogen strengthen

ing at room temperature, so the 304L and 316LN results were similar

223
Although the final fractures were typically in the weld, all specimens

yielded noticeably in the tube material. This indicates that the weld

and tube strengths were well matched at 295 K.

Fracture strengths were substantially (20% or more) above the yield

strengths, which is characteristic of ductile fracture. The variation in

fracture strengths can be correlated to lack-of- fusion defects at the

root of the weld joint. These defects occurred over approximately 10% of

the cross-sectional area and reduced the fracture strengths by as much as 25%. With the exception of specimen 304L-2, significant amounts of plastic extension were observed in both the weld and tube material. For

example, specimen 316LN-4 necked and failed in the tube material at the

center of the gage section. For this reason, the fracture strength of

the tube material was near 700 MPa at room temperature. Overall, the

295 -K test results show that all material combinations deformed and

failed in a ductile manner.

Table 2. Results of the 295—K Tensile Tests

Specimen

Yield

Strength

(MPa)

Fracture

Strength

(MPa)

Plastic

Extension

(mm)

Lack of

Fusion

(%)

316LN-1 490 714 5.4 0

316LN-2 420 651 3.6 0

316LN-3 441 721 7.2 0

316LN-4 448 714 7.1 0

304L-1 434 567 1.4 10

304L-2 420 504 0.6 10

21Cr—6Ni—9Mo1

434 672 4.4 2 21Cr—6Ni—9Mo-2 441 581 1.9 2

B. Tensile Tests at 4 K

Unlike the 295-K specimens, most 4-K specimens failed before reaching the yield-strength criterion. In other words, the plastic extension at fracture (Table 3) was less than that used to define the yield

strength. Factors to be considered with regard to these premature fail- ures are (1) fusion flaws, (2) ferrite content of the welds, (3) inherent

low toughness and notch sensitivity of the 21Cr—6Ni—9Mn alloy at 4 K, 8 and (4) the relative stability of the welds in relation to stress -induced

martensite formation at 4 K.

224
Table 3. Results of the 4—K Tensile Tests

Specimen

Yield

Strength

(MPa)

Fracture

Strength

(MP)

Plastic

Extension

(mm)

Lack of

Fusion

(%)

316LN-5 1484 1820 0.8 0

304L-4 — 1190 0.4 10

304L-5 — 1274 0.0 10

21Cr—6Ni—9Mn-3 — 1050 0.0 2 21Cr—6Ni—9Mn-4 1260 0.0 2

The ferrite of the as -received tube material was measured as 0 FN.

Ferrite measurements, taken on autogenous GTA welds made on the various

stainless steels (including the tube material), ranged between 2 and

3 FN. When dilution effects are taken into consideration, the FN for the

welds between the various material combinations is not expected to exceed

3 FN. Therefore, magnetic responses exceeding 3 FN in deformed regions

of the specimen would imply that some martensite had formed.

Magnetic -response measurements clearly indicated that small amounts

of martensite formed upon deformation of the tube material at 4 K. Results from the deformed weld areas are more uncertain because varying

amounts of plastic deformation occurred in the sample areas of different

welds. The 316LN welds, however were stable: no martensite was detected

after deformation.

Since the initial ferrite contents of the welds were similar, lackof- fusion flaws and low ductility in the welds should have been the major

contributors to premature failure. The effect of any stress - induced mar- tensitic transformations on the failures for these specimens was assumed

to be small.

Upon review of the 4-K test results (Table 3) , the 316LN specimen

had the highest strength and best ductility. Its strength was closest

to that of the tube material (see Appendix A) . Final failure occurred

in the tube material after deformation occurred in both the weld and tube

material. The good behavior of this joint was expected on the basis of

the high ranking given to 316LN in material studies for DOE fusion-energy

applications. 2 However, type 316LN is difficult to obtain, and type 316L

is acceptable for this application.

The 21Cr—6Ni—9Mn specimens failed before reaching the displacement

criterion for the yield strength. The fractures occurred in the welds

and showed some signs of brittleness when examined macroscopically . The

fractures apparently initiated at small lack-of- fusion flaws near the

225
root of the welds. These flaws constituted only about 2% of the weld

area, however, so the result may be more indicative of a low tolerance

for cracks in these welds than a loss of weld cross-sectional area.

The 304L specimens had an average fracture strength 20% higher than

that of the 21Cr—6Ni—9Mn specimens, but they also failed before reaching

a displacement for yield strength. These welds had been produced at a higher heat input, which eliminated the lack-of- fusion flaws, but created

a wider (and weaker) HAZ , the location of their failures.

Notched- Specimen Tensile Tests (Phase 2)

Types 304L and 316L flange specimens were selected for phase 2 testing because they performed better than the 21Cr—6Ni—9Mn specimens in

phase 1 testing. These specimens were notched to simulate a fabrication

defect that might occur in any of four material regions (flange HAZ,

weld, tube HAZ, and tube material). Then they were loaded in tension

to failure.

A. Notched-Specimen Tensile Tests at 295 K

Tensile test results for notched specimens at 295 K (Table 4) indicated that the 316L specimens had higher strengths and greater ductilities than the 304L specimens.

Specimens notched in the flange HAZ . If the welding process did not

degrade the flange materials, specimens notched in the flange HAZ would

be expected to have mechanical properties similar to those of unnotched

specimens (Table 2) because the fracture path is longer on the flange

side. A comparison of the yield strength values of the notched and un- notched 304L and 316L specimens confirms that no severe degradation oc- curred in either of the flange materials. Visual examination indicated

that general yielding initiated in the tube or weld portion of the specimen; no deformation was visible at the flange HAZ notch. Since the

flange HAZ notch was resistant to fracture, the results for the notched

flange HAZ specimens were interpreted as representative of the variation

in the strength and ductility of the weld joint configuration in other

regions. All specimens deformed significantly in the tube material prior

to failure at 295 K, but the 316L specimens were slightly superior to the

304L specimens in both strength and ductility.

Specimens notched in the weld. The largest differences between

strength and plastic extension of the notched specimens occurred in the

weld-notched specimens. The 316L welds had slightly higher yield and

fracture strengths and 38% more plastic extension than the 304L welds.

Specimens notched in the tube HAZ region. The yield strengths,

fracture strengths, and ductilities of the specimens of two flange

materials with notches in the tube HAZ were similar and slightly better

than those of specimens with notches in the unwelded area of the tube.

The similarity in the values for the two HAZ regions was expected since

the welds were produced under the same conditions.

226
Table 4. Results of 295—K Notched- Specimen Tensile Tests

Specimen Notch

Location

Yield

Strength

(MPa)

Fracture

Strength

(MPa)

Plastic

Extension

(mm)

304L-9 Flange HAZ 405 599 2.30

304L-8 Flange HAZ 405 634 3.40

304L-16 Weld 495 714 1.96

304L-15 Weld 515 695 1.80

304L-19 Tube HAZ 493 629 1.36

304L-18 Tube HAZ 500 665 1.64

304L-13 Tube 522 631 1.08

304L-6 Tube 501 625 1.06

316L-3 Flange HAZ 435 656 4.00

316L-10 Flange HAZ 449 673 3.75

316L-7 Weld 558 766 2.38

316L-9 Weld 522 769 2.84

316L-1 Tube HAZ 505 655 1.40

316L-3 Tube HAZ 511 663 1.44

316L-14 Tube 493 626 1.14

316L-17 Tube 497 619 1.08

Specimens notched in the unwelded area of the tube. The pre- sence of a notch increased the yield strength of the tube material by

approximately 16% and reduced the fracture strength by approximately 13%

(cf. the unnotched- specimen values recorded in Table 2). The plasticextension results for these notched- tube specimens indicated that the

tube material had lower ductility than other locations of the specimen

configuration. Comparison of the plastic extension value of the notchedtube specimen (1.1 mm) with that of the unnotched 316LN-4 specimen

(7.9 mm) shows that the reduction in 295-K ductility due to the notch

is appreciable and should be an important fabrication consideration.

B. Notched- Specimen Tensile Tests at 4 K

Table 5 summarizes the results of notched- specimen tensile testing

at 4 K.

Specimens notched in the flange HAZ . Results for specimens notched

in the flange HAZ are representative of unnotched- specimen behavior, as mentioned previously, since failure did not occur at the notch. These

specimens tended to fail at approximately 85% of the tensile strength

determined for the tube (Appendix A). Failure occurred in the tube cross section at the end of the reinforcement weld (Figure 3). At 4 K, the

227
stress- intensity factors for this region of the unnotched- specimen con- figuration resulted in failures at stresses only slightly higher than

those associated with the notched- tube material. Plastic-extension values for the flange HAZ specimens, however, indicated that more ductility

was associated with the flange HAZ failures of unnotched specimens.

Specimens notched in the weld. The notched-weld specimens of the

316L flange material were superior in strength and ductility to notchedweld specimens of 304L flange material, as was the case at 295 K. The

316L notched-weld specimens had slightly higher yield strengths (closer

to the tube yield strengths reported in Appendix A) than the 304 notchedweld specimens. They also had higher fracture strengths than of any of

the specimens with different notch locations and greater ductilities than

the 304L notched-weld specimens.

Specimens notched in the tube HAZ region and specimens notched in

the unwelded area of the tube. Specimens of the two flange materials,

whether notched in the tube HAZ or in the tube, had similar properties at

4 K as well as at 295 K. Three of the yield strength values for the 316L

flange material specimens were undetermined, either because the specimens

broke before our 0 . 05 -mm- extension criterion for yield, or because of a serration in the curve. Nevertheless, the results indicate that welding

did not significantly degrade the tube material at 4 K.

Table 5. Results of 4—K Notched- Specimen Tensile Tests

Specimen Notch

Location

Yield

Strength

(MPa)

Fracture

Strength

(MPa)

Plastic

Extension

(mm)

304L-10 Flange HAZ 1154 1369 0.16

304L-7 Flange HAZ 1125 1408 0.36

304L-17 Weld — 1461 0.20

304L — Weld 1380 1434 0.07

304L-20 Tube HAZ 1307 1414 0.12

304L-21 Tube HAZ 1319 1405 0.08

304L-11 Tube 1281 1387 0.10

304L-12 Tube 1342 1387 0.08

316L-11 Flange HAZ 1254 1519 0.64

316L-12 Flange HAZ 1213 1513 0.72

316L-5 Weld 1446 1682 0.31

316L-8 Weld 1510 1698 0.28

316L-2 Tube HAZ — 1453 0.04

316L-4 Tube HAZ — 1455 0.10

316L-15 Tube — 1383 0.07

316L-16 Tube 1349 1391 0.08

228
Notched- Specimen Fatigue Tests ('Phase 3)

Since phase 2 tests showed that the flange HAZ of notched specimens

was an unlikely fracture path, and the results for specimens with notches

in the tube and tube-HAZ areas were the same for both flange materials,

type 304L and type 316L candidate flange materials could be completely

evaluated in the fatigue test program with only four specimen configurations: one of each of the two flange materials with notches in their

welds, one with a notch in the tube HAZ, and one with a notch in the

tube

.

A set of five specimens was used to determine the fatigue load

for a life near 4 x 10 4 cycles at 295 and 4 K for each specimen configuration. Test results are listed in Tables 6 and 7 and summarized in

Table 8 and Figures 7 through 10. Although the tabular data are more

precise, the figures show the average S—N slope and the data scatter more

clearly. In the figures, the solid and dashed lines were drawn through

the data points to estimate the mean S—N response at 295 and 4 K,

respectively.

Table 6. Results of 295—K Fatigue Tests

Specimen Fatigue Strength

(MPa)

Fatigue Life

(cycles)

304L Weld- 27 412 3.37 X 10 4 304L Weld-28 413 2.56 X 10 4 304L Weld- 29 487 0.94 X 10 4 304L Weld-30 503 1.2 X 10 4 304L Weld-31 412 2.6 X 10 4 316L Weld-22 412 4.8 X 10 4 316L Weld-23 487 1.45 X 10 4 316L Weld-24 487 1.83 X 10 4 316L Weld-25 412 6.21 X 10 4 316L Weld-26 412 4.03 X 10 4 Tube HAZ- 27 412 3.92 X 10 4 Tube HAZ- 28 412 4.5 X 10 4 Tube HAZ- 29 412 5.06 X 10 4 Tube HAZ- 30 487 1.02 X 10 4 Tube HAZ -31 487 1.4 X 10 4 Tube -6 412 2.71 X 10 4 Tube- 18 487 1.01 X 10 4 Tube- 19 487 1 . 17 X 10 4 Tube -20 412 2.98 X 10 4 Tube- 21 412 2.1 X 10 4

229
Table 7. Results of 4—K Fatigue Tests

Specimen Fatigue Strength

(MPa)

Fatigue Life

(cycles)

304L Weld-22 449 1.95 x 10 4 304L Weld-23 487 1.7 x 10 4 304L Weld- 24 487 2.11 x 10 4 304L Weld-25 449 1.7 x 10 4 304L Weld-26 487 1.87 x 10 4 304L Weld-C 389 2.80 x 10 4 304L Weld-D 389 3.75 x 10 4 316L Weld-42 493 1.15 x 10 4 316L Weld-43 493 0.71 x 10 4 316L Weld-44 410 3.15 x 10 4 316L Weld-45 449 1.49 x 10 4 316L Weld-46 389 4.16 x 10 4 Tube HAZ-32 450 2.54 x 10 4 Tube HAZ-33 410 3.10 x 10 4 Tube HAZ-34 410 2.45 x 10 4 Tube HAZ-35 375 3.41 x 10 4 Tube HAZ-36 450 1.40 x 10 4 Tube -37 412 1.16 x 10 4 Tube -38 375 1.27 x 10 4 Tube -39 487 0.63 x IQ 4 Tube -40 450 0.98 x 10 4 Tube -41 487 0.18 x 10 4 Tube -A 340 1.37 x 10 4 Tube-B 290 4.66 x 10 4

Table 8. Estimated Fatigue Strength for

a Fatigue Life of 4 x 10 4 Cycles

Estimated Fatigue Strength (MPa)

Specimen

at 295 K at 4 K

304L Weld 378 357

316L Weld 420 392

Tube HAZ 420 392

Tube 371 287

230
FATIGUE LIFE, cycles

Figure 7. S—N curves for notched-weld specimens

with 304L flange material.

FATIGUE LIFE, cycles

Figure 8. S—N curves for notched-weld specimens

with 316L flange material.

231
STRESS,

MPa

Figure 9. S—N curves for notched- tube -HAZ specimens.

232
A. Notched- specimen fatigue tests at 295 K

Specimens notched in the weld. Examination of the strength at which

the mean S—N lines intersect the design life of 4 x 10 4 cycles (Figures 7 and 8) shows that the type 316L flange specimen with a notch in the weld

had a higher fatigue strength (420 MPa) than that of the 304L flange

specimen (378 MPa) . The scatter in life data is quite small (one standard deviation is nearly 14 MPa) when the complexity of the specimen

design and gage region is considered; therefore, the higher fatigue

strength for the weld to the 316L flange is statistically significant.

The slopes of the S—N curves for the two welds are similar.

Specimens notched in the tube HAZ. The S—N slope for the specimen

notched in the tube HAZ (Figure 9) is similar to that of specimens

notched in the weld (Figures 7 and 8) . Its fatigue strength (420 MPa)

is the same as type 316L notched in the weld. This high fatigue strength

might be attributed to metallurgical changes in the cold-worked tube

structure due to the heat of welding.

Specimens notched in the unwelded area of the tube. The S—N slope

for specimens notched in the unwelded area of the tube is similar to

those for other specimens. The fatigue strength of these specimens

(371 MPa) is nearly the same as that of the type 304L specimen with a notched weld (378 MPa). That the fatigue strength is lower than that of

specimens with notches in the tube HAZ might be attributed to cold working during the formation of the strip or the effects of processing into

tube form.

The 295-K test results in Table 8 can be summarized by stating that

the notched- tube -HAZ specimens and notched-weld specimens with the type

316L flange had the highest fatigue strength for a life of 4 x 10 4 cycles. This is very beneficial because the virtual crack of the tubeto- flange joint ends near both of these materials. The low fatigue

strength of welds to the type 304L flange suggest that this material

should not be used.

B. Notched- specimen fatigue tests at 4 K

The mean S—N responses for specimens tested at 4 K were decidedly

nonlinear, as shown in Figures 7 through 10. This nonlinearity challenges the accuracy of the S—N response determined from a limited test

program. Therefore, more specimens were produced and tested, which

eliminated the necessity to extrapolate the line to 4 x 10 4

. The scatter

in the life data was more at 4 K (perhaps a standard deviation as high as

28 MPa) than at 295 K.

The relative ranking of the estimated strengths of the four specimen

designs was the same at 4 K as at 295 K (see Table 8).

233
O

VP

x 10 4

. “ icrcn

* j «

2

IBT-SZZZ~~B ZZ i 6 r.s - c.er cn .non

.

55 - — _ . V6 — ZS ZZ - • _ S w w-«r.sC £ zz b

253

:< r.o:;.',

.

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_..C _ a. Cffice cf r us cr.

— — •< R. A ov

[c Henry S rrengzh—roughnes

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Oi

P.
4. N. J. Simon and R. ?. Reed, Strength and toughness of AISI 20- and

316 at 4 K, in Materials Studies for Magnetic fusion Zr.err-' Attllcatior.s at Lou Temperatures — IX . R. ?. Reed, ed.. .*2 S IR 86-3051.

National 3ureau of Standards, 3oulder, Colorado, 1936, pp . 2 ~—-1.

5. Standard test methods and definitions for mechanical testing of stee

products — A370, 1 9

3 ~ Book of ASOM Standards . vol. 03. 01. American

Society for Testing and Materials, Philadelphia, Pennsylvania. 193'.

p. 4.

6. D. T. Read and R. L. Tobler, Mechanical property measurements at lov

temperatures , in: Advances in Cryogenic Zntlneerint - Materials

vol. 28, Plenum, New York, 1982, pp. 17—28.

7. Mechanical Testing , vol. 8 of the Ninth Edition of Metals H andbook

American Society for Metals, Metals Park, Ohio, 1933, p. 5~2.

8. R. L. Tobler and D. Meyn, Cleavage - like fracture along slip planes i Pe— 18Cr-3Ni—13Mn—0 . 37N austenitic stainless steel at liquid helium

temperature, in: Materials Studies for Mametlc fusion Zr.err-' Applications at Low Temperatures — XI . N3SIR 88-3032. R. ed.

,

National Bureau of Standards, 3oulder, Colorado, 198

pp. 89-101.

Iti

*u
APPENDIX A. MECHANICAL PROPERTIES OF THE 21Cr-6Ni-9Mn TUBE AT 4 K

This adjunct study was prompted by (1) lack of information on the

mechanical properties of 21Cr—6Ni—9Mn tube at 4 K and (2) information

received during the testing program that seamless tube with a 1.2-mm wall

thickness would replace the welded tube with a 0.9 -mm wall thickness from

which our test specimens had been constructed.

To evaluate the differences in tube material, both tubes were tested

in uniaxial tension at 4 K. A simple reduced-section specimen design

with a 6 -mm width and a 40 -mm gage length by tube thickness was used.

To be conservative, the loading direction selected was normal to the

tube length because the tube is weaker in this direction than along its

length. This loading direction also enabled the seam weld to be cen- tered in the gage length and transverse to the loading direction of the

specimen.

The values for yield and tensile strength listed in Table Al show

no weakening of the specimen due to the weld. The higher yield strength

for the welded specimen is attributed to the data scatter caused by our inability to completely remove the existing bend from the tube during

specimen preparation without affecting the mechanical test results.

Therefore, the yield strength values are nominal and should not be used

in design calculations.

Although the elongation values (based on a nominal 25 -mm gage

length) indicate that the two tube thicknesses could be used interchangeably in the study of flange materials, we chose to use only the

tube with 0.9 -mm- thick walls to minimize the variables in the study. For

the same reason, we excluded seam welds from the gage lengths of these

specimens

.

Table Al. Mechanical Properties for the 2lCr—6Ni—9Mn Tube at 4 K

Specimen

Yield

Strength*

(MPa)

Tensile

Strength

(MPa)

Elongation

(%)

seamless — 1.2-mm wall 1420 1841 2 . 6

1

welded — 0.9-mm wall

only base metal

in gage length

1246 1806 2. it

welded — 0.9-mm wall

transverse seam weld

in gage length

1425 1806 17. 5*

* For information only. Since the tubes -could not be totally

straightened without introducing permanent deformation, the

yield strength data contain substantial scatter.

t Broke outside the nominal 25-mm gage length.

^ Broke within the gage length.

236
STRENGTH AND TOUGHNZSS CE AN AllTCOift?/- -DEOXIDIZE!

ECR CRYOC-INIC SERVICE^

C. N. MeCowan and T. A. S lever

t

National Bureau of Standards

Boulder , Colorado

D. J. Kotecki

Teledyne McKay Research Laboratory

York, ? eimsyIvani a

The strength and toughness of an aluminum deoxidized stainless steel veld (Fe—13Cr—13Ni—2QMnr-0. 17C) were evaluated for

cryogenic service. At 4 K, the yield and tensile strengths of

the veld were approximately 850 and 950 MPa, respectively; its

fracture toughness [Kj c (J)] was 85 MPa'm’1

. At 76 K, its Chart”

V-notch absorbed energy was 42 J. The strength—toughness relationship of this weld at cryogenic temperatures was no better

than that of more conventional, silicon- deoxidized welds. Targe

(2- to 5-nm), aluminum- rich inclusions were observed in the

weld, and data related to the role of inclusions in the ductile

fracture process are reported.

I . INTRODUCTION

The fracture toughness of AISI type 316LN stainless steel at cryogenic temperatures is affected by inclusion content, size, and distribution. 1 Two parameters controlling the inclusion morphology are the

welding process and the deoxidation system of the welding electrode.

Silicon is the principal deoxidant in stainless steel electrodes . but

manganese and chromium often combine with the Si0 2 because together a more stable compound is formed. The result -is a closely spaced dispersion of spherical MnSi03 inclusions with 0.1- to 1-^m diameters.

The 4-K fracture toughness of welds is lower than that of matching

composition base materials. 2 Preliminary studies indicate that the ::: nary cause for this lower toughness is the higher inclusion content ir. the welds. 3 Since weld toughness limits the design of cryogenic structures, other inclusion- control processes and deoxidation systems are being sought.

Tnis paper was prepared specifically for this report: it will - published elsewhere.

237
This study investigates the use of an aluminum-based deoxidation

system for stainless steel welds as a means to improve the toughness at

4 K. To avoid precipitation of AiN particles, the weld was strengthened

by carbon and manganese additions rather than nitrogen. To increase the

toughness, levels of chromium and nickel were chosen to produce a weld

with low ferrite and high nickel content. 4 ’ 5

II. MATERIALS AND PROCEDURE

A. Weld Material

The weld composition and ferrite content are given in Table 1. To

obtain the desired weld composition, a thin strip of type 304L stainless

steel was used as the sheath for the flux core electrode. The flux com- ponents and the remainder of the alloy were added to the center. The

electrode produced was 2.4 mm (3/32 in) in diameter.

Table 1. Weld Composition (wt.%) and Ferrite Number (FN)

c Mn Si Cr Ni N Al 0 P S FN

0.172 20.0 0.33 i-l CM VO 13.4 0.05 2.0 0.001 0.023 0.003 0.3

B. Welding

Welds were made using a flux core arc (FCA) electrode. A 25 -mm- thick mild- steel plate material was beveled to produce a 45° included

angle for the single -vee weld joint configuration. The base material was double buttered (surfaced with the welding electrode to prevent dilution

effects) in accordance with AWS A5.22-80 prior to making the weld. After

buttering, a root gap of 12.5 mm was used to ensure that an ample volume

of undiluted weld metal was available for test-specimen preparation. The

welds were made at 275 A and 28 V with a travel speed of about 380 to

400 mm/min (15 to 16 in/min)

.

C. Mechanical Testing

Uniaxial tensile specimens with a 12. 5 -mm diameter (conforming to

ASTM A 370) were used for 293-K testing. Specimens with a 6 . 25 -mm diameter were used in tests at 4 K. These all-weld-metal tensile specimens

were removed from the welds so that the tensile axis was parallel to

the welding direction. One 293-K tensile test and two 4-K tests were

performed. Specifics of the liquid-helium tensile- testing apparatus are described elsewhere. 6

238
The 25 . 4-mm- thick compact tension specimen was side-grooved to

enhance plane strain conditions at the crack front; the notch was parallel to the welding direction and oriented to result in crack growth along

the welding direction. Fracture toughness was measured at 4 K by using

the J-integral test method designated by ASTM E 813-81 and the singlespecimen unloading-compliance method. 7 The Kj c (J) value was calculated

from

KIc (J) = (Ji c-E) h (1)

For this calculation, the value used for Young's modulus (E) at 4 K was

200 GPa.

III. RESULTS AND DISCUSSION

A. Metallography

The weld had a cellular solidification structure and appeared to be

fully austenitic (Figure 1). Inclusion diameters in the weld ranged from

greater than 5 pm to less than 0.5 pm. The large, angular inclusions

were rich in aluminum, according to energy-dispersive x-ray analysis of

the polished surface with a scanning electron microscope. Manganese -rich

inclusions were also identified in the matrix with x-ray analysis, but

they occurred much less frequently than the aluminum-rich inclusions,

and they were generally smaller (less than 1 pm) with a more spherical,

glassy appearance.

Fig. 1. Fully austenitic weld with large, faceted,

aluminum-rich inclusions.

239
B. Tensile Testing

The tensile test results (Table 2) show that the yield strength

of the weld increased from 383 to 850 MPa when the test temperature was decreased from 293 to 4 K. The tensile strength increased from 565 MPa

at 293 K to 950 MPa at 4 K.

% Table 2. Tensile Test Results

Temperature

(K)

Yield Strength

(MPa)

Ultimate Strength

(MPa)

Elongation

(%)

293 383 565 41

4 849 1000 20

4 849 924 12

On the basis of results of two previous studies, the relative

strengthening contributions of alloying elements in the weld were estimated. The strengthening terms taken from these studies were

a. 14 MPa/wt.% for manganese, from a study of type 18Cr—8Ni stainless steel welds that contained up to 10 wt.% manganese; 8

b. 1700 MPa/wt.% for carbon, from a study of the strengthening

contributions of carbon and nitrogen in type 300 series stainless steels. 9 Using these terms, we calculated that carbon would contribute 292 MPa,

and manganese would contribute 280 MPa to the yield strength of the weld

at 4 K. Subtracting these contribution^ from the total yield strength of

850 MPa leaves a base strength of 278 MPa, which includes the strengthening due to other solid-solution alloying elements in the weld. This re- sult agrees well with the 275- to 300-MPa constants used in the equations

from which the carbon and manganese coefficients were obtained. Although

this excellent agreement may be somewhat fortuitous, the comparison indicates that the substitutional strengthening of manganese in the weld may

be comparable to the interstitial strengthening of carbon.

The substitutional strengthening of aluminum was not considered

in the estimates of strength. Bulk x-ray analysis, however, indicated

that the weld matrix probably had more than 1 wt.% aluminum in solution.

Since the volume occupied by aluminum atoms at substitutional lattice

positions is greater than that of manganese atoms, the misfit and lattice

distortion is greater, and some degree of strengthening by aluminum would

be expected. 10 ' 12 Rough estimates of atomic size indicate that aluminum

240
additions would result in a lattice expansion more similar to that of

molybdenum than that of manganese. 13 Strengthening coefficients reported

for molybdenum at 4 K range from 30 to 80 MPa/wt . %

.

1

4

'

1

7

C. Fracture Testing

Charpy V-notch test results for five weld specimens were an average

absorbed-energy value of 41.6 J (30.5 ft-lbf) and a lateral expansion of

0.41 mm (0.016 in) at 76 K (Table 3). These values are similar to the

40-J and 0.45-mm values of 18Cr—8Ni stainless steel SMA welds strengthened to 850 MPa at 4 K by nitrogen and manganese additions, but they are somewhat lower than the 55 -J and 0.75 -mm values of the SMA welds with

14 wt.% nickel. 15

Table 3. Charpy V-Notch Test Results at 76 K

Sample Absorbed Energy

(J)

Lateral Expansion

(mm)

1 45.2 0.46

2 41.2 0.41

3 38.2 0.41

4 44.1 0.48

5 38.2 0.33

The fracture toughness of the weld was 85 MPa -in 2 at 4 K. The

Kj c (J) calculated for the weld is comparable to the lower limit of

the strength—toughness trend observed for type 308L and 316L stainless

steel weld metal at 4 K. 2 The upper band limit, at a 850-MPa strength

level, however, is near 150 MPa-m , and base metal values are even

higher. Therefore, the strength—toughness relationship of the carbonstrengthened, aluminum- deoxidized weld composition evaluated in this

study was no better than that of the more conventional nitrogenstrengthened stainless steel welds.

D. Fractography and Inclusion Characterization

In a continuing effort toward improved modeling of the effect of

inclusions on crack initiation and growth for ductile fracture processes,

several kinds of measurements are a routine part of our micros true tural

characterization procedure. Data gathered for the weld composition investigated in this study include inclusion counts on a mechanically polished surface, ductile-dimple counts on fractured surfaces, and diameter:,

for inclusion—dimple pairs found on the fracture surfaces.

241
LiulW/l* V.^V .Ss'TfV.v

‘J,- r

' :•!& • ‘•'VsfiS *8p;jiRS!

^ \

*

' ^

•*>

Vt« ^ *

' * <-*. ». ' 4

+ -

- PNISV* •«-••

)ws«sg&

Fig. 2. Ductile -dimple morphology observed near the notch

of a 4-K compact- tension specimen.

The fracture surfaces of the test specimens showed no signs of brittle behavior during failure (Figure 2) . Ductile-dimple morphologies in areas near the notch on the compact tension specimens and in areas near

the center of the tensile specimens were similar in appearance. The

faceted appearance of the larger aluminum- rich inclusions is shown in

Figure 3

.

When cross sections taken normal to the fracture surface of 4-K

tensile specimens were examined, void growth at inclusion sites was limited to a thin layer adjacent to the fracture surface or cracks that

had branched away from the fracture surface. As shown in Figure 4,

little or no void growth was observed at inclusions even when they were located very near the crack surfaces. At distances greater than 50 /xm

from the fracture surface, the inclusion—matrix interface generally

appeared to be coherent, but several areas contained microcracks. We

speculate that these areas may be due to microvoid coalescence at inclusion clusters or to weld flaws that later combine and cause fracture.

Similar microcrack formation and linking processes have been postulated

as a failure mechanism for more conventionally deoxidized, nitrogenstrengthened welds with MnSi0 3 inclusions. 1 The weld examined in this

study, however, had fewer voids at inclusion sites than welds with MnSi03 inclusions. So, there may be significant differences in the critical

stress and strain for void nucleation of the two inclusion types.

242
Fig.

Fig. 4.

3. Fractograph of a 4-K tensile specimen showing the large,

faceted, aluminum-rich inclusions present in the weld.

Voids located only at inclusions near cracks (crack branch)

in cross sections normal to the fracture surface

.

243
Inclusion counts from polished weld samples indicate an inclusion

density of approximately 7,500/mm2

, which is low compared to the 30,000

to 40,000/mm2 for flux core arc (FCA) , shielded metal arc (SMA) , and gas

metal arc (GMA) welds evaluated in previous studies. 1 The low inclusion

density reflects the low oxygen content of the weld metal. The planar

inclusion spacing (proportional to the inverse square root of the inclusion density 18

), therefore, is larger than those of previously evaluated

welds. Since the frequency of large inclusions observed on the fracture

surfaces was high, any beneficial effect of the large inclusion spacing

(11.5 /im) on the toughness of this weld composition appears to have been

offset by the increased average size of the inclusions. Adverse effects

on the toughness due to aluminum can not be ruled out, however; reductions in Charpy V-notch absorbed energies have been correlated to aluminum contents in ferritic steels. 19

The ductile-dimple densities estimated from the fractured tensile

and compact tension specimens had similar values—44,500 dimples/mm2

. The ratio of planar inclusion densities to dimple densities is approximately 6, lower than the values of 8 to 11 previously found for welds. 3 In general, these ratios may serve as a macroscopic roughness parameter:

as dimple densities approach planar inclusion densities, toughness de- creases. A ratio of 2 would indicate equal densities for these quantities, since a correction factor of 2 is generally used for the 2-D to 3-D

surface conversion. 20 For now, too few data are available to evaluate

the use of this variable in fracture models.

The inclusion- to-dimple diameter ratio is used to compare the ductile flow of the matrix during the fracture process of welds. The 0.55

ratio for the weld of this study is higher than the 0.25 to 0.35 ratios

for previously evaluated welds, implying that there was less local plastic flow before fracture. 3

IV. CONCLUSIONS

1. The aluminum-based deoxidation system used for the FCA weld evaluated

in this study did not improve the fracture toughness of stainless

steel welds at 4 K.

2. Inclusions found in the aluminum- deoxidized welds were large and appeared to be faceted crystals rather than glassy spheres, which are observed in silicon-based slag systems.

3. The use of aluminum deoxidation systems in stainless steel welds may

result in increased toughness if the inclusion size can be substantially reduced.

ACKNOWLEDGMENTS

This work was supported by the U.S. Department of Energy, the Office

of Fusion Energy.

244
REFERENCES

1. C. N. McCowan and T. A. Siewert, Inclusions and fracture toughness

in stainless steel welds at 4 K, Advances in Cryogenic Engineering -

Materials, vol. 34, Plenum Press, New York, 1988, pp . 355—342.

2. R. L. Tobler, T. A. Siewert, and H. I. McHenry, Strength-toughness

relationship for austenitic stainless steel welds at 4 K, Cryogenics, vol. 26, 1986, pp . 392—296.

3. T. A. Siewert and C. N. McCowan, The role of inclusions in the fracture of austenitic stainless steel welds at 4 K, Welding Metallurgy

of Structural Steels, ed. J.Y. Koo, The Metallurgical Society, New

York, 1987, pp. 415—428.

4. E. R. Szumachowski and H. F. Reid, Cryogenic toughness of SMA aus- tenitic stainless steel weld metals: Part 1 — Role of ferrite,

Welding Journal, vol. 57, no. 11, pp .. 325s—333s.

5. R. P. Reed, P. T. Purtscher, and K. A. Yushchenko, Nickel and nitrogen alloying effects on the strength and toughness of austenitic

stainless steels at 4 K, Advances in Cryogenic Engineering - Materials, vol. 32, Plenum Press, New York, 1986, pp. 43—50.

6. D. T. Read and R. L. Tobler, Mechanical property measurements at

low temperatures, Advances in Cryogenic Engineering — Materials,

vol. 28, Plenum Press, New York, 1982, pp . 17-28.

7. R. L. Tobler, D. T. Read, and R. P. Reed, Strength/toughness relationship for interstitially strengthened AISI 304 stainless steel

at 4 K temperature, Fracture Mechanics, ASTM STP 743, American

Society for Testing and Materials, Philadelphia, Pennsylvania, 1981,

pp. 250-268.

8. C. N. McCowan, T. A. Siewert, R. P. Reed, and F. B. Lake, Manganese

and nitrogen in stainless steel SMA welds for cryogenic service,

Welding Journal, vol. 66, no. 3, 1987, pp . 84s-92s.

9. R. P. Reed and N. J. Simon, Low temperature strengthening of austenitic stainless steels with nitrogen and carbon, Advances in Cryogenic Engineering - Materials, vol. 30, Plenum Press, New York,

1984, pp. 127-136.

10. P. Haasen, Solution hardening in f.c.c. metals, Dislocations in

Solids, ed. F. R. N. Nabarro, North-Holland, New York, 1976,

pp. 155-189.

11. F. R. N. Nabarro, The theory of solution hardening, Philosophical

Magazine, vol. 35, 1977, pp . 613-622.

12. R. Labusch, Statistische Theorien der Mischkristallhartung , Acta

Metallurgica, vol. 20, 1972, pp . 917-927.

245
13. B. D. Cullity, Elements of X-Ray Diffraction, 2nd Ed., Addison

Wesley, Reading, Massachusetts, 1978.

14. R. P. Reed, Austenitic stainless steels with emphasis on strength

at low temperatures, Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures — XI, NBSIR 88-3082, National

Bureau of Standards, Boulder, Colorado, 1988, pp. 1—37.

15. C. N. Mccowan and T. A. Siewert, Influence of molybdenum on the

strength and toughness of stainless steel welds for cryogenic

service, Welding Metallurgy of Structural Steels, ed J.Y. Koo , The

Metallurgical Society of the AIME, New York, 1987, pp . 427-438.

16. 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. 30, Plenum Press, New York, 1986, pp . 57—64.

17. R. Miura, Development of high strength structural alloys for

superconducting reactor, Japan Steel Works, Hokkaido, Japan;

unpublished.

18. E. E. Underwood, Quantitative Stereology, Addison-Wesley , Reading,

Massachusetts, 1970.

19. G. E. Linnert, Welding metallurgy: carbon and alloy steels, vol. 2,

American Welding Society, Miami, Florida, 1967.

20. E. E. Underwood, Estimating feature characteristics by quantitative

fractography , Journal of Metals, April 1986, pp. 30—32.

246
INCLUSIONS AND FRACTURE TOUGHNESS IN STAINLESS STEEL WELDS AT 4 K*

C. N. McCowan and T. A. Siewert

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

Eight stainless steel welds were studied to evaluate the

effect of inclusion content on the fracture toughness at 4 K.

The loss in toughness associated with the inclusion spacing

agreed with the trend previously determined between inclusion

spacing and toughness in stainless steel base materials. This

finding indicates that the smaller inclusion spacings observed

in welds are the principal reason why welds have lower fracture

toughness than base materials of similar compositions at 4 K

and why some welds are tougher than others. Indus ion- size

distributions were found to be similar for the welds. The

oxygen content of the welds was evaluated as an alternate means

of representing the effect of inclusions on the weld toughness

at 4 K.

INTRODUCTION

Type 316LN stainless steel is emerging as the preferred structural

material for superconducting magnet casings designed to operate at 4 K.

It has a higher strength than the older 316L grade and comparable tough- ness at cryogenic temperatures. 1 ' 3 Unfortunately, welding electrodes t matching this composition and strength have substantially poorer tough- ness at 4 K than the stainless steel plate. 4

This study investigates the effect of inclusions on toughness as a . step toward explaining the lower strength—toughness relationship that is* observed for welds.

Expansion of a previous study of three welds to eight welds. The

earlier study was reported in Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures — X. This paper was published

in Advances in Cryogenic Engineering — Materials , vol. 36, Plenum,

New York, 1988, pp . 335-342.

247
Although the role of inclusions in the dimple rupture fracture pro- cess has long been recognized, attempts to model effects of inclusions

on the fracture toughness have only recently been undertaken. 5

* 7 Traditionally, the process of dimple rupture is separated into three stages:

(1) void nucleation, (2) void growth, and (3) void coalescence. Nonmetallic inclusions are a common site for void nucleation. However, the

relative amounts of the energy measured in the fracture toughness test

due to void nucleation as opposed to void growth and coalescence may

vary, depending on the inclusion morphology and composition, the testing

conditions, and the strength of the matrix material. 6

EXPERIMENTAL PROCEDURES

Eight stainless steel welds were selected for evaluation in this

study. Only welds having FN values of less than 3 were selected to avoid

any appreciable effect on toughness due to this variable. The compositions and ferrite numbers (FN) of the welds are given in Table 1.

The test matrix includes welds with varying nonmetallic inclusion

content (a result of different welding processes), strength, and toughness. The welds are characterized and referenced in Table 2.

Table 1. Compositions (wt.%) and Ferrite Content (FN) of the Welds

Weld c Mn Si P S Cr Ni Mo N 0 FN

497-5 0.,01 2.2 0 .,46 0.,010 0 .,013 18 15 2. 3 0 ,03 0 ,,059 0 .9

508-2 0 .,02 1.4 0.,47 0.,020 0..003 18 12 2. 1 0 ,.07 0.,025 3 .0

1616 16 16 2 0 ..040 0

1 0 ,.03 1.7 0..58 0,.033 0 ..011 19 14 2. 3 0 .04 0 ,.070 0 .1

5 0 .06 1.7 0..56 0 ..033 0..009 19 13 2. 3 0 .04 0,.063 1 .2

10 0 .03 5.4 0 .26 0 ,.006 0 ,.007 18 20 0 . 10 0 .16 0 ,.037 0

11 0 .03 5.4 0 .26 0 .006 0 ,.007 18 20 0 . 10 0 .16 0 ,.005 0

12 0 .03 5.4 0 .26 0 .006 0 ,.007 18 20 0 . 10 0 .16 0 ,.005 0

248
Table 2. The 4-K Mechanical Properties, Inclusion Data,

and Welding Information for the Welds

Weld Ref

.

4-K Yield

Strength

(MPa) (MPaVm)

Inclusion

Density

(no . /mm2

)

Mean

Inclusion

Diameter

( 4m)

Welding

Information

497-5 8 608 159 33,173 0.41 GMAW, tubular

electrode

,

99Ar—10 2 508-2 8 779 163 24,983 0.29 GMAW, 99Ar-10 2 1616 9 607 235 19,050 0.47 FCAW

1 10 694 179 31,796 0.39 SMAW

5 10 924 143 37,587 0.38 SMAW

10 11 1015 203 11,317 0.45 GMAW, 98Ar—20 2 11 991 287 9,801 — GMAW, Ar

12 1027 263 8,300 — GMAW, Ar

Welds 11 and 12 were made using the same GMA electrode and welding

configuration as weld 10. 11 The shielding gas used for these welds,

however, was commercial- grade argon instead of the 2%-oxygen, 98%-argon

mixture used for weld 10. In addition, a second shielding enclosure was fabricated to reduce the oxygen contents around the arc. For weld 11,

a droplet transfer was established (20 V; 175 A; 12-mm tip-to-work distance), and for weld 12, spray transfer was established (25 V; 280 A;

18 -mm tip-to-work distance).

The oxygen contents of the welds were determined with an interstitial element analyzer. Sections (0.5 to 1 g) were cut from test speci- mens using an abrasive cut-off saw and then cleaned in hydrochloric acid

prior to analysis. The average oxygen content of five samples was re- ported. Standard deviations did not exceed 0.01 wt.%.

Inclusion measurements were made on mechanically polished weld

samples taken from the fractured tensile specimens. The data for welds

10, 1616, 297-5, 508-2, 1, and 5 were gathered using an image analysis

program on a scanning- electron microscope (SEM) at a magnification of

3000X. Back- scattered electrons were used for imaging. Three hundred

inclusions were counted per sample. Inclusion-density data for welds 11

and 12 were obtained visually at 3000X. These welds were recently

tested, and size distribution data has not yet been accumulated. The

average inclusion spacing (A) was calculated as the inverse square roo*:

of the inclusion density.

249
ANT DISCISSION

Distribution ar.d Density of Inclusions

Three general types of inclusion- size distributions were observed

for the welds. The distribution curves for the welds having inclusion

densities larger than 30,000/mm2 are similar and can be characterized by

one lir?e in Fig. 1. As might be expected, the two S>LA welds are included

in this group. The number of small- diameter inclusions in these welds

increased rapidly up to a peak at approximately 0.25 pm. A similar

inclusion- size distribution was found for weld 497-5. This weld was made using a tubular GMA electrode and a 99% - argon , 1%-oxvgen shielding

gas mixture. The inclusion density of the FCA weld (1616) was only

19,000, mm2

; the distribution curve was similar in shape to that of the

GMA weld made with a 2%-oxygen, 9S%-argon shielding gas. 10 These two

welds are characterized by a second line; they show a broader distribution of inclusion sizes through the peak region. The two types of curves discussed thus far do not differ greatly; the majority of inclusions in

these welds have diameters within the range 0.15 to 0.45 pm, the mean

diameter values differ by less than 0.07 pm, and the curves have "tails"

as larger inclusion diameters are approached. The only significant

difference was the distribution for weld 508-2 (a GMA weld made using a 1%-oxvgen, 99%-argon shielding gas). More than 50% of the inclusions in

this weld had diameters of less than 0.3 pm.

Results here indicate that, with the exception of weld 508-2, the

principal difference among the inclusions in these welds is the inclusion

density and not the distribution. Veld 508-2, however, clearly indicates

that different types of size-distribution curves are possible.

Figure 1. Inclusion- size distribution curves for the welds.

250

* 10

3

Figure 2. The Quality Index (normalized to 12 wt.% nickel) at 4 K,

versus the inclusion spacing for type 316 stainless steel

base materials. Scatter bands represent 2 SDs. 13

e

•

CM

(0 a.

X

0) T3_C

>.

ra3

o

X103

Figure 3. The Quality Index (normalized to 12 wt.% nickel) at 4 K

versus the inclusion spacing for the welds.

252
between inclusion spacing and the Quality Index at 4 K agrees with that

found for the base materials. In fact, the magnitude of the effect that

inclusions have on the fracture toughness would indicate that this

parameter alone may explain the lower strength—toughness relationship

observed for the welds.

Oxvgen Content as a Measure of Toughness

The oxygen content of a weld deposit is another means by which the

inclusion content can be measured. When the bulk of the weld-metal

inclusions are assumed to be oxides, the inclusion volume fraction is

proportional to the oxygen content of the weld deposit. 14 On this basis,

oxygen content was compared to weld toughness in Fig. 4. The Quality

Index, normalized to 12 wt.% nickel, was used to indicate more clearly

any relationship between the oxygen content and the toughness. The

Kj,i(d)-^y values decreased as the oxygen content increased.

Oxygen, wt. %

Figure 4. The Quality Index (normalized to 12 wt.% nickel)

versus the oxygen content of the welds.

The usefulness of oxygen content as a measure of the inclusion con- tent is a promising, but rather simplistic approach. It requires a similar distribution of the inclusion volume fraction for the welds, and it

does not include the influence on toughness of the more fundamental quantity, inclusion spacing. Oxygen content, when used in conjunction with

inclusion distribution data, may provide more meaningful parameters to be

used in the prediction of fracture toughness. The inclusion spacing can

be defined in terms of the inclusion volume fraction: 15

A = cf" 1 / 3 ?,

where c is a constant, f is the volume fraction, and r is the average inclusion radius. In this study, the use of [0 2 ] -1//3 -r and other complex

variables did not improve the correlations between the fracture toughness

253
and oxygen content. However, measurement of oxygen content is easy and

reproducible, and with further study of its interaction with inclusion

size distributions, an alternative means to express the effect of inclusions on toughness may be developed.

Fracture Mechanisms

To obtain a more accurate model of the influence of inclusions on

the fracture toughness of stainless steel welds at 4 K, the relative

importance of void nucleation versus void growth in the formation of an internal crack must be understood. As a first step, metallographic

sections near the centers of two tensile specimens tested at 4 K were

examined.

Voids formed at both the small and large inclusions in weld 10

(Fig. 5a). Only a small amount of void growth occurred, however, during

the 26% uniform elongation in this specimen. There are some indications

of void growth just before impingement and coalescence in Fig. 5a (upper

right). Void coalescence, however, was rarely observed in the specimen.

When it was observed, commonly it had occurred very near the fracture

surface. The more traditional void impingement occurring in these areas may be the result of accelerated void growth due to the triaxial stresses

imposed during the passage of the crack front. 6 In this case, coales- cence would be the result of a crack rather than a source for crack formation. A comparison of void sizes on the polished surface with those on

the fracture surface illustrates the sudden increase in void size that occurs during fracture.

Figure 5. Cross-sectional views along the axis of tensile specimens

fractured at 4 K: void impingement in weld 10 (a)

;

void

linkage in weld 497-5 (b)

.

254
Evidence that crack formation at 4 K was occurring by the linking of

small voids rather than by void impingement was observed in weld 497-5.

The crack perpendicular to the applied stress in Fig. 5b appears to be

forming by "microcrack- linking" (lower, center region of the figure).

Small cracks can be seen joining several of the voids in this area.

In summary, the observations of fractured tensile specimens indicate

that void growth plays a lesser role in the fracture of stainless steel

welds at 4 K than expected. Similar findings have been reported in

stainless steel base materials at 4 K. 16 Since the extent of void growth

is small compared with the inclusion spacing, it appears that the stress

or strain required for the onset of void growth and a factor to increase

the energy required for fracture as the inclusion spacing increased would

be two meaningful parameters for modeling the fracture event.

CONCLUSIONS

1. The inclusion spacing was the principal reason for the lowering of

the fracture toughness for stainless steel welds below the values

observed in the base materials at 4 K.

2. The inclusion density was highest for the SMA welds and decreased

for the GMA as a function of the shielding gas and type of electrode

used.

3. The distribution of inclusion sizes in the welds did not differ

substantially on the basis of the welding process used.

4. The oxygen content of the welds is another way to express the effect

of inclusions on toughness.

ACKNOWLEDGMENT

This study was sponsored, in part, by the Office of Fusion Energy,

U.S. Department of Energy.

REFERENCES

1. G. M. Goodwin, Weld. J . 64:19-25 (1985).

2. S. Yamamoto, N. Yamagami, and C. Ouichi , in: "Advances in Cryogenic

Engineering—Materials," vol. 32, Plenum Press, New York (1985),

pp. 57—64.

3. N. J. Simon and R. P. Reed, J . Nucl . Mater . : 141—143: 44—48

(1986).

4. R. L. Tobler , T. A. Siewert, and H. I. McHenry, Cryogenics 26:392—

396 (1986).

5. G. Leroy, J. D. Embry, G. Edwards, and M. F. Ashby, Acta Metal 1 . . 29:1509-1522 (1981).

6. R. H. Van Stone, T. B. Cox, J. R. Low, Jr., and J. A. Psioda, Int . Met. Rev. 30:157-179 (1985).

255
10 ,

11

12

14,

15

16

.rtr. and. r . n . -roes, Xe:;.-

:icple and D. J. Kotecki, in:

mans . 8A: 1165-1176 (1977).

"Material Studies for Magnetic

Fusion Energy Applications

H

C.

N.

R.

K.

M.

P.

t Low Temperatures — IV" (N3SIR 81-

1545), National Bureau of Standards, Boulder, Colorado 1981,

pp. 303-322.

I. McHenry and T. A. Whipple, in: "Material Studies for Magnetic

Fusion Energy Applications at Low Temperatures — III" (N3STR 80-

1529), 3ouIder, Colorado, 1980, pp . 155—166.

T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas Jr.,

in; "Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures — II" (NBSIR 79-1609), National Bureau of

Standards, Boulder, Colorado, 1979, pp . 315—352.

N. McCowan, T. A. Siewert, and R. L. Tobler, J. Er.e. Mater.

Techno 1. 108:340-343 (1986).

J. Simon and R. ?. Reed, in "Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures — XI" (NBSIR 88-

3082), National 3ureau of Standards, Boulder, Colorado, 1988,

pp. 71—87.

?. Reed, N. J. Simon, P. T. Purtscher, and R. L. Tobler, in:

"Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures — IX" (N3SIR 86-3050), 1986, pp . 15—26.

Easterling, in: "Introduction to the Physical Metallurgy of

Welding," Butterworths , Boston (1983), p. 85.

F. Ashby and R. Ebeling, Trans. AI ME 236:1396 (1966).

T. Purtscher, R. P. Walsh, and R. P. Reed, in: "Materials Studies

for Magnetic Fusion Energy Applications at Low Temperatures — XI"

(N3SIR 88-3082), National Bureau of Standards, Boulder, Colorado,

1988, pp. 113-129.

256
HOHMETALUCS

257

NORMETALUCS PROGRAM

CONTRIBUTOR: H. Ledbetter

OBJECTIVES

• Determination of composite-material physical properties that are useful in engineering design, properties such as Young's modulus,

torsion modulus, and internal friction.

• Development of analytical and theoretical models to describe and

predict these properties.

• Study of physical property interrelationships.

RESEARCH HIGHLIGHTS

• For glass-fiber epoxy-matrix composites, the ambient- temperature

torsional modulus, G, and the associated internal friction, Q

-1

, were

measured. To understand how G depends on fiber volume fraction,

existing models were modified to allow for nonhomogeneous fiber distribution. All simple models fail to explain the dependence of the

torsion-mode Q

-1 on fiber volume fraction.

• For glass-fiber epoxy-matrix composites, the torsion modulus and the

mechanical- loss spectrum, Q

-1 (T), were measured between 295 and 4 K.

Mechanical loss arose entirely from the matrix material, which re- sponded like a typical polymer. Fiber-matrix interfaces showed zero internal friction. Thus, the glass-epoxy composites showed lower

damping than the epoxy.

259

INTERNAL- FRICTION SPECTRUM OF A GLASS—EPOXY COMPOSITE*

M. Weller

Max -Planck- Ins ti tut fur Metallforschung

Ins ti tut fur Werkstoffwissenschaften

Stuttgart, Federal Republic of Germany

H. Ledbetter

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

Using a computer-controlled inverted torsion pendulum

at frequencies near 1 Hz , we determined the internal -friction

spectrum of a uniaxially fiber-reinforced composite. The com- posite comprised glass fibers in an epoxv-resin matrix. We

studied three fiber contents: 0, 41, and 49 volume percent.

Three internal-friction peaks appeared: above 300 K, near

200 K, and near 130 K. We label these a, >9, and 7 , respectively. They correspond closely to a, and 7 peaks found

previously in many polymers. We failed to see an internalfriction peak for either the glass or the glass—resin interface. Between 300 and 4 K, the torsion modulus increased in

the resin by a factor of 3.39 and in the 0.49 glass -epoxy by

a factor of 1 . 68 .

INTRODUCTION

In composites, internal friction arises from three sources: the

occluded phase, the matrix phase, and the interface. In structural

composites, the reinforcing phase usually possesses high elastic

stiffness and low internal friction. Thus, composites with good lowinternal- friction) interfaces possess a lower internal friction than

the matrix phase. How low depends on the reinforcement-phase volume

fraction.

*Intended for Journal of Applied Physics.

261
Whether usual composite -material interfaces contribute measurably

to internal friction remains uncertain. Some studies report an interface contribution. 1 Others found none. 2

' 4

This question assumes importance for several reasons. First, internal friction is a fundamental physical property whose understanding

pays many dividends. 5 Second, internal friction (or damping) enters many

composite -material engineering applications. 6 Third, we need to know

whether internal friction provides a useful tool for studying interfaces?

A recent review on composite -material plastic deformation concluded that

"interfaces play the most important role in the behavior of a composite." 7 Fourth, current theoretical studies 8 ’ 9 on interface damping would

profit enormously from experimental studies.

This study used a torsion pendulum to determine the low-frequency

internal- friction spectrum of glass-fiber-reinforced epoxy-matrix com- posites. Previously, for these composites, Ledbetter et al

.

4 used kilohertz standing-wave methods to measure the Young's modulus and internal

friction.

100 pm

Fig. 1. Photomicrograph of a composite transverse to uniaxial glassfiber bundles in an epoxy-resin matrix. Fiber volume, c, is

0.41; fiber diameter equals 7 fim. Nonhomogeneous fiber distribution results from fiber bundles used in manufacture.

262
EXPERIMENT

Materials

Composites were made at NBS by using vacuum- impregnation methods

described by Kasen. 10 The manufacturing method yielded 3 -mm- diameter

cylinders up to 30 cm long. Except for cutting to length, we measured

these as received. The matrix material consisted of a commercial epoxide

resin formulated for use in a low- temperature radiation environment. The

reinforcing phase consisted of tows of commercial glass fibers 7 pm in

diameter with a reported mass density of 2.60 g/cm3 and a Young's modulus

of 72.4 GPa. Figure 1 shows a photomicrograph of the 41% glass composite. Table 1 contains the mass densities of all three materials studied.

Fiber volume fraction was determined from three parameters: glass mass,

glass mass density, and mold volume.

Table 1. Measured properties of glass—epoxy composites

c

P (g/cm3

) G(kHz)

(GPa)

G(Hz) G 0 s t

(K)

•k t

(K)

0.0 1.115 1.49 1.0 6.65 35.21 652 397

0.41 1.654 5.41 3.6 4.50 52.40 957 463

0.49 1.852 5.25 3.8 5.35 66.22 956 621

c - fiber volume, p - mass density, G - axial torsion modulus,

G 0 - torsion modulus at 0 K, s - an adjustable parameter re- lated to zero-point vibrations; T - temperature; t and t* are parameters relating to the effective Einstein temperature. G(Hz)

are relative values. G 0 and s represent values normalized by

dividing by G(295).

Measurements

The cylindrical specimens served as the suspension of an inverted

Ke-type torsion pendulum. The pendulum was excited to an initial sur- face strain of approximately 1.5-10" 5

. The near- 1 -Hz free-decay torsion

amplitude was recorded with a light-beam system and a semiconductor

detector. Temperature was measured with a calibrated thermocouple with

an inaccuracy of 0.05 K. The overall systematic temperature error is

estimated to be 0.5 K. Inside the pendulum a static pressure of 0.1 kPa

was maintained; by a cryopump, the partial pressure of other residual

gases, such as oxygen and nitrogen, was reduced to approximately 0.1 Pa.

For temperature measurements, which were controlled by a desk computer,

the mean heating rate was 1.5 K/min. By using a digital measuring system, three quantities—frequency, f; temperature, T; and internal friction, Q' 1 —were calculated immediately, punched on tape, and plotted.

263
A Dewar enabled the specimen to be cooled to near liquid-helium temperature . With an internal electrical heater, the specimen could be heated

to approximately 500 K. Because of high damping, the heating was stopped

at approximately 330 K.

Fig. 2. Variation of f

2 and Q" 1 for three volume fractions

versus temperature.

RESULTS

Figure 2 shows the principal results: internal friction and squared

frequency versus temperature.

The squared frequency is proportional to the axial torsion modulus,

G; ii

G - 4pf2 i

2 (1)

where p denotes mass density and i, specimen length. The f

2 -versus -T

behavior is regular. One can fit such regular behavior to many empirical

264
equations. One such equation, based on an Einstein- oscillator tcde..

originated with Varshni : 12

G - G0 — s/[exp(t/T) - 1] (2)

where G 0 denotes the zero- temperature torsion modulus; s, an adjustable

parameter related physically to zero-point vibrations; and t, an adjustable parameter related to the effective Einstein temperature . Table 1 shows the fitting results. Fitting Eq. (2) over the 4- to 340-K region

gave poor results. Parameters in Table 1 represent the 140- to 340-K

region, except t 2 , which represents the 4- to 340-K region.

All three materials show qualitatively similar Q* 1 (T) behavior: a background that decreases with decreasing temperature and three internal

friction peaks 1 near 130 K, near 200 K, and a above 300 K.

DISCUSSION

The present Q

_

1 (T) results resemble closely those reported previously for polymers. For example, for nylon, Wert and Weller 13 reported

internal-friction peaks at 130, 200, 290, and 340 K. With small temperature shifts, these a, 0 , and 7 peaks occur in various polymeric solids

that include oil shale, asphalt, wood, cork, amber, and coal . 14

Wert and Weller 14 proposed mechanisms for these three relaxation

peaks. Gamma arises from twisting under the mechanical stress of linear

cross-link methylene units: —CH 2 —CH 2 —CH2 — . Beta arises from reorientation under stress of a small unit such as OH, CH 2 0H, or COOH. Alpha

arises either from a large reorientation of main chains or from a glass

transition.

That all three observed peaks arise from the polymeric matrix means

that, in the temperature region studied, neither the reinforcement phase

nor the interface contributes measurably to internal friction. Absence

of an interface Q'

et al . 4 agrees with the kilohertz measurements of Ledbetter

A theoretical analysis of interface internal friction by Datta et

al . 8 ’ 9 shows that interface damping increases with increasing vibrational

frequency, increasing reinforcement-phase volume fraction, increasing oc- clusion size, increasing occlusion elastic stiffness, decreasing matrixphase elastic stiffness, and increasingly spherical occlusion shape. Increasing interface thickness nay increase or decrease internal friction

S waves attenuate more than ? waves. Thus, we need further studies on a wider range of composites before settling the question of internal friction in composite interfaces.

Finally, we comment on the failure of the f

2 measurements to fit

Eq. (2) over the entire temperature range. The good fit obtained between

265
140 and 340 K suggests a low- temperature change of state. We can imagine

this either as an elastic stiffening occurring during cooling below 140 K

or an elastic softening occurring above 140 K during cooling.

CONCLUSIONS

From this study of a glass-fiber-reinforced, epoxy-resin composite

between 4 and 340 K, there emerged five conclusions:

1. Three internal -friction peaks were discovered by torsion-pendulum

measurements

.

2. 'All three peaks correspond to those that occur in macromolecular

solids, such as nylon.

3. Neither the glass phase nor the glass—resin interface showed an internal -friction peak.

4. Determination of the significance of internal friction in compositematerial interfaces requires more measurements.

5. The shape of the f

2 -versus-T curve suggests a different state below

approximately 140 K.

ACKNOWLEDGMENTS

Dr. M. B. Kasen of NBS produced the composite materials. OwensComing (Granville, Ohio) furnished the glass fibers. S. A. Kim of NBS

assisted with numerical analysis. This study was sponsored, in part, by

the Office of Fusion Energy, U.S. Department of Energy.

REFERENCES

1. K. Shimizu, in Composite Materials. Proceedings of Japan—United

States Conference (Japanese Society of Composite Materials, Tokyo,

1981), p. 111.

2. N. Igata and A. Kohyama, Trans. Jap. Soc. Compos. Mater . 2, 30

(1976).

3. H. M. Ledbetter, J. Phvs. (Paris) 46, CIO-573 (1985).

4. H. M. Ledbetter, Ming Lei, and M. W. Austin, J . App 1 . Phvs . 59, 1972

(1986).

5. N. G. McCrum, B. E. Read, and G. Williams, Anelastic and Dielectric

Effects in Polymeric Solids (Wiley, London, 1967).

6. "Vibration Damping, 1984 Workshop Proceedings," Air Force Wright

Aeronautical Laboratories Report AFWAL-TR-84-3064 (November 1984).

7. Y. Flom and R. J. Arsenault, J. Met . 38, no. 7, 31 (1986).

8. S. K. Datta and H. M. Ledbetter, in Mechanics of Material Interfaces

(Elsevier, Amsterdam, 1986), pp . 131—142.

9. S. K. Datta. H. M. Ledbetter, Y. Shindo , and A. H. Shah, in Review

of Progress in Quantitative Nondestructive Evaluation , vo 1 . 6

(Plenum, New York, 1987), pp . 1075—1084.

266
10 . M. B. Kasen, in Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures — VIII . NBSIR 85-3075, National

Bureau of Standards, Boulder, Colorado (1985), p. 87.

11. H. F. Pollard, Sound Waves in Solids (Pion, London, 1977), p. 89.

12. Y. P. Varshni, Phvs . Rev. B 2, 3952 (1970).

13. C. A. Wert and M. Weller, J. AppI. Phvs . 53, 6505 (1982).

14. C. A. Wert and M. Weller, J. Phvs. (Paris) 46, CIO-561 (1985).

267

TORSION MODULUS AND INTERNAL FRICTION OF A FIBER-REINFORCED COMPOSITE*

H. Ledbetter and Ming Lei

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

By a kilohertz - frequency resonance method, we determined

the torsion modulus and internal friction of a uniaxially

fiber-reinforced composite. The composite comprised glass

fibers in an epoxy-resin matrix. We studied three fiber contents: 0, 41, and 49 vol.%. The internal friction failed

to fit a classical free -damped- oscillator model where one as- sumes a linear rule -of- mixture for three quantities: oscillator

mass, force constant, and mechanical-resistance constant. The

torsion modulus fit approximately a plane-wave-scattering

ensemble -average model. The microstructure showed strong non- homogeneity of fiber distribution. Considering this nonhomogeneity yielded better agreement between model and observation.

Thus, torsion-modulus measurements provide a method to detect

and quantify the nonhomogeneity of fiber distribution.

Previously, the authors 1 reported measurements and modeling of the

Young's modulus and associated internal friction of a uniaxial-fiberreinforced epoxy-matrix composite. Here we report new measurements and

modeling of torsion modulus and internal friction.

The materials used for the present study are the same as those used

for the previous Young' s- modulus study. 1 Again, we studied three fiber

contents: 0, 41, and 49 vol%. The experimental methods were described

previously. 1 To get the torsion modulus, T, we used the relationship

T - pv2 - 4pf|iJ. (1 )

*Intended for Journal of Applied Physics.

269
Here, p and v denote mass density and torsion-wave sound velocity;

SL s denotes specimen length; fs denotes specimen resonance frequency.

For torsion-mode oscillation, we obtained fs from the following

relationship

m

q

f

2 - f2 + (f2 - f

2

) — s o o q ms

( 2 )

Here, subscripts o, q, and s denote complete oscillator, quartz - crystal

,

and specimen, respectively; f denotes resonance frequency, m denotes

mass, and d denotes diameter. The specimen's internal friction, Qg 1

, is

determined by the following relationship:

m. q :

1 + <q:‘

o o q:

1

) — • me (3)

Table I contains the properties of constituent materials. Table II

shows the principal measurement results. Figures 1 and 2 show graphs of

T and Q" 1 versus fiber volume fraction, c.

Table I. Properties of Constituent Materials

P (g/cm3

)

E

(GPa)

A* (GPa)

B

(GPa)

k

(GPa)

A

(GPa)

V

Glass 2.577 75.7 30.3 50.5 60.6 30.3 0.250

Resin 1.155 3.11 1.17 3.11 3.50 2.34 0.333

Table II. Measured Properties of Glass -Epoxy Composites

c t 3 (GPa)

Q" 1 (10- 3

) P

(g/cm2

)

fs (kHz)

^s

(cm)

0.0 1.75 77.62 1.155 41.20 1.495

0.41 5.41 40.5 1.654 40.48 2.235

0.49 5.25 27.02 1.852 40.756 2.067

270
Fig. 1. Torsion modulus versus fiber volume fraction. Points represent measurements. Lower curve represents homogeneous- fiber

distribution case. Upper curve represents an extremely

nonhomogeneous case: a single fiber. Middle curve represents

a model calculation that allows for fiber bundles . Dashed

line represents a linear rule-of-mixture

.

Fig. 2. Internal friction versus fiber volume fraction. Points represent measurements: squares, kHz; circles, near 1 Hz. Curve

represents a classical free-damped oscillator model, which

fails to explain Q

_1 (c).

271
First, we consider a model for the axial torsion modulus, T 3

. From

standard sources, 2 we know that the torsion modulus along principal axes,

i, can be expressed in terms of the elastic-compliance tensor,

T.

1

2

s ijij + ^ikik

where i ^ j ^ k. Thus, along the x x axis

% 2 2

'12 12 + S 1313 * 6 6 + s 55

(4)

(5)

We applied the usual Voigt contraction scheme to the second expression in

Eq. (5). Similarly,

and

>2121 + S 2 323 8 6 6 +

S 3 i3i + S 3 2 3 2 S55 + s 44

( 6 )

(7)

We consider the continuous -fiber -reinforced composite as a transversely isotropic material. In this case, we have S 44 - S 55 . Therefore,

Eq. (7) becomes

t 3 ( 8 )

Here, denotes a Lame constant (shear modulus). Subscript LT denotes

shear deformation in any plane containing the x3 axis, that is, the fiber

axis. Using energy methods, Hashin and Rosen3 derived relationships for

the effective moduli of a continuous -fiber -reinforced composite in which

the fibers are distributed randomly and homogeneously. They obtained the

following relationship for the shear modulus:

2c (/if - /%)/%

MLT “ Aim +

/if + /% - c(/if - /%)

(9)

Here, c denotes fiber volume fraction; subscipts m and f denote matrix

and fiber. Using the elastic constants given in Table I, we calculated

/iLT and show the results in Fig. 1. The model calculation gives results

well below measurements. We believe this disagreement arises from non- homogeneous fiber distribution, which is visible in Fig. 3. To deal with

the nonhomogeneity, we modified the above model to contain two steps. In

the first step, we apply the effective -moduli relationships to uniaxial

glass -fiber bundles. By assuming fiber volume fraction, cff,, equal to

0.87 in the bundles, we obtain the bundle properties. In the second

step, by using the bundle properties as new "fiber" properties and taking

bundle volume fraction c^, instead of c, in Eq. (9), we obtain the non- homogeneous composite property. The volume fraction c^ relates simply

to c:

272
( 10 ) c °b ” c fb

Figure 1 shows that the effective shear modulus for nonhomogeneous fiber

distribution agrees well with measurements, but we found that this approach fails for high volume fraction, above approximately 0.60 for the

present system. Figure 1 also shows the torsion modulus of a clad

composite cylinder calculated by the relationship 4

T - pv2 [cpf + (1 - c)pm ][^m - (/% - /if)c 2

] Pm ~

(Pm “ Pf)° 2 ( 11 )

This represents an extremely nonhomogeneous case: a single glass fiber

surrounded by a cylindrical tube of epoxy-resin matrix.

100 urn Fig. 3. Photomicrograph of studied composite transverse to uniaxial

glass-fiber bundles in an epoxy matrix. Fiber diameter equals

7 pm. Nonhomogenous fiber distribution results from fiber

bundles used in manufacture.

273
Second, we consider internal friction Q" 1

. In the previous study, 1 starting with an equation for free vibrations and a linear rule-ofmixture for three quantities—mass, m; force constant, k; and mechanicalresistance constant, b—we gave a relationship for Q"

1

:

Q

'

1 Qm

1 f

1 + rh. • sf)

kf

+

r c 1 2 Pf r c

i

3 Pf ^ k +

m

[l - cj Pm

L-i - cj P k

m m J Here

,

bm

(mmkm

)^

( 12 )

(13)

For torsion-mode vibration, we use the following relationship for the

force constant:

k - Ta . (14)

Here, a denotes a characteristic interatomic spacing. This led to kf =

9.78 N/m and 1% - 0.117 N/m. We treated am as the single adjustable fitting parameter in Eq. (12), using am — 1 A. Figure 2 shows the model

calculation and kilohertz measurements for internal friction, Q” 1

. Fig- ure 2 also shows Q” 1 determined by the torsion pendulum at a frequency

near 1 Hz. 5 Unlike the internal friction measured in longitudinal -mode

vibration, the model calculation predicts results much lower than the

measurements. Therefore, this model for internal friction fails to apply

to the present case. The reason for the failure requires further study.

ACKNOWLEDGMENT

This study was sponsored, in part, by the Office of Fusion Energy,

U.S. Department of Energy.

REFERENCES

1. H. M. Ledbetter, Ming Lei, and M. W. Austin, J. Appl. Phys . 59, 1972

(1986).

2. R. F. S. Hearmon, Applied Anisotropic Elasticity (Oxford University

Press, London, 1961), p. 51.

3. Z. Hashin and B. W. Rosen, J. Appl. Mech. 31, 223 (1964).

4. A. E. Armenakas , J. Acoust. Soc. Am. 30, 439 (1965).

5. H. Ledbetter and M. Weller, Materials Studies for Magnetic Fusion

Energy Applications — XI, NBSIR 88-3082 (National Bureau of Standards, Boulder, Colorado), p. 261.

274
TECHNOLOGY TRANSFER

275

TECHNOLOGy TRANSFER PROGRAM

LEADER: N. J. Simon

CONTRIBUTORS: L. Delgado, C. J. King, R. P. Reed

OBJECTIVES

• 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

• Preparation of an annual report, as well as monthly highlight reports

to the Office of Fusion Energy, U.S. Department of Energy

• 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

• The Tenth Annual Cryogenic Structural Materials Workshop was organized; it was held October 8—9, 1987 in Pleasanton, California.

• Materials Studies for Magnetic Fusion Energy Applications — X, (NBSIR

86-3067, 444 pages, 1987, editor R. P. Reed) was prepared, published,

and distributed.

• Additional 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, creep, fatigue, and thermal properties has been com- pleted; data have been extracted and analyzed. Evaluation of documents for similar properties of C17000—C17510 beryllium copper has

also been completed; a large amount of data on tensile properties has

been extracted and analyzed.

• Handbook pages and supporting documentation covering tensile, electromagnetic, elastic, fatigue, and thermal properties of C10100—C10700

copper have been distributed. Recent NBS data on creep in C10400

copper are being analyzed to elucidate creep mechanisms at cryogenic

temperatures. Handbook pages on C17000—C17510 beryllium copper tensile and electromagnetic properties are in preparation.

277

TENTH ANNUAL

CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

Sheraton Pleasanton

Pleasanton, California

October 8—9, 1987

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 1987

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

• New low- temperature materials developments in the

United States and international research cooperation

79

10th annual

Cryogenic Structural Materials Workshop

Program

THURSDAY, OCTOBER 8

8:30 a.m. Introduction V. Der, OFE

COMPACT IGNITION TOKAMAKS

8:45 a.m. Current Status of the CIT G. Brown, PPPL

9:30 a.m. Structural Problems of the Alcator C-MOD H. Becker, MIT

10:00 a.m. Compression Tests of Copper -Inconel

Laminates for CIT Magnets

R. Fields, NBS

10:30 a.m. Mechanical Properties of Electron Beam

Welds in Thick Copper

T. Siewert, NBS

11:00 a.m. Effects of Grain Size and Cold Rolling

on Cryogenic Properties of Copper

R. Reed, NBS

11:30 a.m. Reinforced Copper Composites for the

Alcator C-MOD

H. Becker, MIT

STRUCTURAL ALLOYS FOR LARGE-SCALE FUSIONI. AUSTENITIC STEELSENERGY DEVICES -

1:30 p .m. Cryogenic Magnet Case and Distributed

Structural Materials for High-Field

Superconducting Magnets

L. Summers,

LLNL

2:00 p . m. Low Temperature Properties of HighManganese Stainless Steels

R. Reed, NBS

2:30 p .m. The Influence of High Magnetic Fields

on the Mechanical Properties of

Austenitic Steels

J.

U.

Morris , Jr .

,

California

3:00 p .m. Effect of Void Nucleation on Fracture

Toughness

P. Purtscher,

NBS

3:30 p .m. Fatigue Crack Growth in Metastable

Austenitic Steels

Z.U.

Mei,

California

281
THURSDAY, OCTOBER 8, continued

4:00 p . m

.

Inclusions and Fracture Toughness in C. McCowan,

Stainless Steel Welds at 4 K NBS

5:30 -

7:30 p .m. Reception

FRIDAY, OCTOBER 9

STRUCTURAL ALLOYS FOR LARGE-SCALE FUSION ENERGY DEVICES - II

8:30 a.m. Status of U.S.—Japan Development R. Tobler, NBS

of Test Standards

9:00 a.m. Welding of High -Manganese Cryogenic Steels J. Chan,

U. California

9:30 a.m. The Cryogenic Mechanical Properties of J. Morris, Jr.

Structural Aluminum Alloys U. California

10:00 a.m. Structural Properties of Nb 3 Sn M. Strum, LLNL

Superconducting Wires

10:30 a.m.

STRUCTURAL COMPOSITE INSULATORS

Composite Insulation Materials: Requirements L. Summers,

for a Small Superconducting Ignited Tokamak LLNL

11 :00 a.m. The Low-Temperature Neutron Irradiation H. Kerchner,

Facility ORNL

11:30 a.m. Review of Gamma and Neutron Irradiation M. Kasen,

Effects on Organic Composite Degradation CTD

Cryogenic Temperatures

12:00 p . m. SSC Dipole Magnet Suspension System Design T. Nicol,

Fermi lab

1:30 p .m.

STRUCTURAL COMPOSITE INSULATORS: LLNL REQUIREMENTS

Open Discussion: Radiation-Tolerant Insu- Discussion

lating Materials for Next Generation Leader:

Fusion Machines V. Der, OFE

282
TENTH ANNUAL CRYOGENIC STRUCTURAL MATERIALS WORKSHOP —

SUMMARY OF TECHNICAL PRESENTATIONS

COMPACT IGNITION TOKAMAKS

. The Compact Ignition Tokamak

ignition reactor, will operate at

coils. Toroidal field coils will

up stresses in the nose region of

yield strength of the cold-worked

(CIT) prototype, a near-term fusion

77 K and use copper or copper alloy

produce a field of about 16 T, setting

400 to 550 MPa, considerably above the

C10700 copper conductor.

Reviewing this and other materials problems of the CIT, G. 3rown

(PPPL) detailed the use of explosively bonded Inconel 718 to strengthen

the conductor structure. The nose region is one of the most critical

areas in the design, but the effects of high radiation levels and high

compression loads on the insulation are also of concern.

H. Becker (MIT) described the structural problems of the Alcator

C-MOD, a 9-T tokamak with reinforced copper rectangular field coils.

Fatigue survival controls most of the structural design, since the

tokamak is expected to undergo 50,000 operational cycles from 77 K to

near ambient temperature.

R. J. Fields (NBS) reported results of recent static and cyclic

compression tests of specimens of a copper—Inconel laminate for the CIT.

The 0.2% -offset yield strength of the laminate was about 550 MPa; cyclic

tests revealed some creep at higher loads.

The design of the CIT requires joining segments of 25 -mm- thick

C10700 copper. T. Siewert (NBS) reported mechanical properties at 77

and 295 K of electron beam weldments of this material. Weld yield

strength was only about 30% of the base metal strength in reducedsection specimens.

R. P. Reed (NBS) presented results of his studies on strengthening

C10400 copper by cold working and reduction of grain size. The best

ratio of strength to resistivity at cryogenic temperatures was obtained

with material that was cold-rolled about 10%.

H. Becker (MIT) described results of compression tests on stacks

of loose copper and steel plates (to simulate explosively bonded C1C~T1

copper—Inconel 718 composites) . Preliminary results indicated that

copper annealed to a yield strength of 48 MPa could survive a -07 -MPa

load.

283
STRUCTURAL ALLOYS FOR LARGE-SCALE FUSION ENERGY DEVICES -

I . AUSTENITIC STEELS

The prototype fusion reactor that will follow the CIT is expected to use superconducting magnets that operate near 4 K. L. T. Summers (LLNL)

discussed the structural requirements of cable- in- conduit conductors for

the prototype TIBER- II (Tokamak Ignition/Burn Experimental Reactor) and

ITER (International Thermonuclear Experimental Reactor) . He reported a proposal to eliminate fatigue problems for the conduit -conductor system

by keeping all stresses compressive in the center toroidal field coil

set

.

R. P. Reed (NBS) characterized recently developed high -manganese

austenitic steels with regard to strength and toughness mechanisms. He

discussed the physical properties, such as thermal expansion, that are affected by the antiferromagnetic transitions in these steels.

J. W. Morris, Jr. (University of California—Berkeley) described

recent measurements of the effects of high magnetic fields on fracture

toughness of a metastable austenitic steel. Earlier studies at the

University of California - Berkeley had found that fatigue crack growth

rate was unchanged in an 8-T field, so the fracture toughness results

were unexpected.

P. Purtscher (NBS) presented evidence that void nucleation (rather

than void growth) is the controlling mechanism in cryogenic fracture

toughness of 304- type alloys.

Z. Mei (University of Califomia-Berkeley) reported effects of the

austenite to martensite transformation on fatigue crack growth behavior

of 304L and 304LN austenitic steels.

Type 316LN stainless steel is also a candidate structural alloy for

superconducting magnet casings. Because the weld toughness of 316LN and

other alloys at 4 K generally is lower than that of the base alloy,

stainless steel welds of varying inclusion and oxygen content were fabricated with several different welding processes in work carried out by

C. N. McCowan (NBS). Like base-metal toughness, the weld-metal toughness

at 4 K increased with reduced inclusions and oxygen content. Since welds

have more inclusions than base metals, the differences in the inclusion

spacing may explain differences in weld- and base-metal toughness.

STRUCTURAL ALLOYS FOR LARGE-SCALE FUSION ENERGY DEVICES - II

R. T. Tobler (NBS) described recent round-robin tests on 4-K tensile

and fracture properties carried out in the United States and Japan. New

standards for 4-K tensile and fracture- toughness testing are the outcome

of this program.

284
J. W. Chan (LBL) discussed 4-K fracture toughness of a highmanganese austenitic steel, 18Mn-16Cr-5Ni-0 . 2N. Fracture toughness

was related to 5 -ferrite microstructure in the weld.

In contrast to the fracture toughness of austenitic steels, that

of the aluminum-lithium alloy 2090 increases at cryogenic temperatures.

J. Glazer (LBL) reviewed mechanical properties of this alloy and discussed a strain-hardening model that accounts for the observations.

In a new test facility that allows testing in fields up to 15 T,

M. J. Strum (LLNL) recently studied the mechanical properties of Nb 3 Sn

superconducting wires. He reported that the conductor—stabilizer com- posite is subject to failure at low strains in or near the operating

range

.

STRUCTURAL COMPOSITE INSULATORS

L. T. Summers (LLNL) descussed the insulation requirements for new

prototypes. The inner turns of the toroidal field coils of the TIBER- II

compact tokamak are expected to experience a high neutron fluence of

10 19 n/cm2

, because the reduced size is achieved with less nuclear

shielding. The tum-to-tum insulation is expected to consist of ceramic

or glass fibers impregnated with an organic resin that will be introduced

after the heat treatment of the conductor. The lack of an adequate data

base for these insulators at radiation levels of 10 10 Gy is a serious

problem.

H. R. Kerchner described the testing facilities at the Low-Temperature Neutron Irradiation Facility (LTNIF) , which are designed to irradiate and test materials at 4 K. When fully operational, this facility

will provide needed measurements of radiation effects on cryogenic

materials for fusion reactors. The fast neutron flux (energy >0.15 Mev)

will be 1.4 x 10 17 n/m2 s.

M. B. Kasen (CTD) reviewed gamma and neutron irradiation effects on organic composites. The most comprehensive recent studies were carried

out by S. Egusa at ANL on composites fabricated from several types of

epoxy and one polymide with type E glass or graphite reinforcements.

T. H. Nicol (FNAL) described the composite suspension system for the

proposed SSC dipole magnets. Both graphite and fiberglass-reinforced

composite tubes will be utilized.

STRUCTURAL COMPOSITE INSULATORS : LLNL REQUIREMENTS

V. Der (OFE) led the discussion of insulator requirements for the

next generation of fusion reactors. The discussion objective was to

identify development and testing needs for insulators and to recommend a base-program action plan to address design issues relating to insulators

285
for both ITER and CIT. Requirements were set forth for an ideal database

on composite- insulator tests in a neutron fluence and spectrum equivalent

to that of TIBER- II. The CIT requirements for compression strength of

insulators were reiterated. New methods of surface preparation to improve interfacial chemical bonding between matrix and reinforcement were

discussed (failure in composite insulators often occurs at the interface) . Several particpants expressed the need for data on 4-K irradiation of insulators followed by testing without warm-up.

286
UNITED STATES- JAPAN

DEVELOPMENT Of TEST STANDARDS

287

UNITED STATES-JAPAN DEVELOPMENT OF TEST STANDARDS PROGRAM

PRINCIPAL INVESTIGATOR: R. L. Tobler

CONTRIBUTING ORGANIZATIONS: National Bureau of Standards,

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.

• Five drafts of a proposed standard 4-K tensile test procedure were

written and submitted for review at Japanese domestic and United States—

Japan workshops. The fifth revision has been formally submitted to ASTM.

• Five drafts of a proposed standard for 4-K fracture toughness testing

were written and reviewed at Japanese domestic workshops. The fifth

revision was submitted to 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 is scheduled for presentation at the 1988

United States—Japan International Workshop.

• An experimental research program was completed by participants at NBS 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. Seven papers describing the

results have been published.

289
Five 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

« At ASTM technical meetings, the draft standards were presented to the

task groups that will process these standards. The meeting on fracture

testing was held in Reno, Nevada on April 16—17, 1988; the meeting on tensile testing, in Baltimore, Maryland on May 16—17, 1988.

290
May 11, 1988

MEMORANDUM FOR: Dr. S. Shimamoto

Japan Atomic Energy Research Institute

and

Dr . H . I . McHenry

Fracture and Deformation Division

National Bureau of Standards

FROM: Ralph Tobler

Fracture and Deformation Division

National Bureau of Standards

SUBJECT: Trip Report — ASTM Technical Committee Meetings

- Reno/Sparks, Nevada, April 26—27, 1988

- Baltimore, Maryland, May 16—17, 1988

INTRODUCTION

With ASTM, we are establishing standards for tensile and fracture

testing at 4 K. Special standards are necessary because of unusual ma- terial behavior at this temperature. The standards we wrote as part of

the Japanese cooperative program are now being reviewed in ASTM subcommittees. I participated in selected meetings to answer questions, obtain guidance, and assimulate new advances in room temperature test

methodologies

.

FRACTURE MEETINGS IN RENO

1. Crack Tip Opening Displacement (CTOD) Concepts (E24.05).

A new CTOD test will soon be officially established by ASTM. At

this meeting John Underwood persuaded the attendees that the value of

the rotational factor, r

p

, should be modified, but all other negative

returns from a recent ballot were resolved. This standard should pass on the next ballot. It provides an alternative method of toughness

evaluation for metals of intermediate toughness when Kj c and Jj c tests

are not valid. This standard could be used for cryogenic testing, and

its development is influencing work in other subcommittees, as noted

below.

291
2. Ductile- to-Brittle Transitions (E24.08).

The meeting on ductile- to-brittle transitions included presentations

on analysis routines (Ted Anderson) , measurement techniques (George

Irwin) and design philosophy (Mike Dawes). The basic question is: how

do you run tests in the transition region? To resolve the major problem,

which is high scatter in measurements, the Metals Properties Council

(MPC) plans to run a data-analysis round robin in tlieir Japanese Cooperative Program; they are currently searching for a large quantity of wellcharacterized material. CTOD tests are applicable to this program.

3. Initiation Concepts (E24.04).

The Initiation Concepts group just revised the Jj c standard E813-87

and published it in the 1987 Annual Book of Standards. The most significant change is the updated definition of Jj c

. The meeting consisted of an introduction, a preview of new directions for future work, and two

technical presentations : Summary of Suggested Changes for Testing in the

Transition Region (J. Joyce) and Update on Standardization of Fracture

Testing at 4 K (R. Tobler) . Twenty experts were present: Drs . Joyce and

Hackett (the chairmen), Vessel, Merkle, Ernst, Landes, Druce, and others.

The new directions for work are

1. J- testing in the transition region

2. Dynamic loading effects

3. New specimen geometries, especially the arc tension specimen

4. Small -specimen testing

5 . Cryogenic testing

The predominant theme now is transition-region characterization.

The thinking in this group seems to reflect developments in the CTOD

group. As a result, a new parameter, J c , may evolve to characterize

those nonlinear materials that fail before developing a full J^-curve.

For example, J c may be defined for abrupt fractures occurring after

stable crack extensions of not more than 0.15 mm. (We already specify

such a parameter for use in our proposed 4-K standard)

.

An interesting difference in material behavior is this: slow, stable

ductile cracking occurs at room temperature, whereas fast unstable ductile cracking occurs at 4 K, but the latter is overlooked in the existing

ASTM standard (see ASTM E 813-87, section 4.3).

Although cryogenic behavior is not a "hot topic" for many of these

specialists, several helpful suggestions were made ’in the discussion

following my presentation. In particular, Hugo Ernst (Georgia Tech)

wants to cooperate with us on the subject of tearing instability theory

as it relates to the material behavior at 4 K. John Merkle recommended

that a task group be formally established for this work, including other

people with cryogenic experience. The task group assignment is now being

pursued by the subcommittee chairmen.

292
MEETINGS IN BALTIMORE

4. Temperature Effects (E28.10)

This meeting with eleven attendees was chaired by R. McDemus

.

The agenda included sign-up and introduction of new members, approval

of previous minutes, task group reports, and new business.

Task Groups : Task Group E28. 10.02 is home for our 4-K tensile standard. At the present time, this standard is in ballot at the subcommittee level. The outcome is forthcoming June 3. At this meeting,

four ballots were returned with highly favorable comments. At least

eight more ballots are required. If no negatives are sustained, this

standard will then advance to general committee ballot and finally to

society ballot.

Four other task groups dealing with high- temperature methods were

reviewed: standard practices for creep (E139) , time to rupture (E292)

,

compression (E209) , and tension (E21) . Documentation for multiplespecimen test apparatus is needed. In this regard, I can contribute

to these task groups by visiting the National Research Institute for

Metals in Tokyo to obtain information on such apparatus that are used

in Japan.

New Business : The possibility of a general low- temperature tensile

test standard was discussed. Since E28 started in 1984, there has

never been one. Consequently, Thomas Wilmink agreed to head a new task

group designated E28. 10.06 — Low- Temperature Tensile Testing. The plan

is as follows: First, our 4-K tensile standard will be processed. After that, E28.10.06 will extend the scope of that standard to include

all low temperatures. NBS was offered an active role in this.

5. Committee E28 Orientation

This informative introductory meeting gave me a general overview

of the scope of the main committee and the standardization procedures

of ASTM.

The structure of the main committee, E28, was described, the

chairman of the various subcommittees were introduced, and the procedures for creating a standard were reviewed.

Today there are nineteen active subcommittees in E28. The majority of standards within their jurisdiction are undergoing revisions

necessitated by three powerful forces: computer automation, international unification of standard practice, and national accreditation of

laboratories

.

293
CONCLUSIONS

New developments in toughness testing are evolving. Appropriate

methodologies are at hand for all types of material behavior: linear

elastic (Kj c ) , elastic—plastic (CTOD, or J c ) , and fully plastic (Jj c ).

In the future, these methods in refined states may be integrated into

a single comprehensive standard.

Our 4-K tensile standard is now well along, undergoing due process.

The fracture standard is not quite as far along; the following course

seems recommendable : review it with the Japanese in forthcoming meetings

(May 26—31) create the sixth draft with all desired changes, form 'a new

task group in E24.04, and initiate the balloting procedure.

294
LOADING-RATE EFFECTS ON DISCONTINUOUS DEFORMATION

IN LOAD -CONTROL TENSILE TESTS*

T. Ogata and K. Ishikawa

Superconducting and Cryogenic Materials Research Group

National Research Institute for Metals

Ibaraki, Japan

R. P. Reed and R. P. Walsh

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

In load-control tensile tests at liquid helium temperature, an abrupt and large discontinuous deformation occurs,

which differs from the discontinuous deformation obtained from

displacement-control tests. We investigated the effects of

loading rate, varied from 0.5 to 5000 N/s, on the tensile properties of AISI 304L, 310, and 316LN steels at 4 K. A large deformation, near 40% strain, occurred in AISI 310. At the high

loading rates, the ultimate strength of these materials was 65%

of the strength obtained in displacement-control tests; the

initiation strength of discontinuous deformation was also less.

INTRODUCTION

The discontinuous deformation that occurs during tensile tests

at temperatures near 4 K is caused mainly by the extremely low specific

heat and thermal conductivity of the material at this temperature. In

load-control tensile tests, the discontinuous deformation is large and

abrupt. 1 ’ 2 This deformation differs from that reported in displacement - control tests, 3

' 7 because the amount of displacement in load-control

tests is not limited by the testing machine, and no load drop occurs during the deformation. Discontinuous deformation in load-control tests

results in a lower ultimate strength than that measured with displacementcontrol tests. So, it is important to clarify these deformation behaviors

and to be able to predict the mechanical properties when a nonrestricted

deformation force is applied to structural materials at low temperatures

In this study, we investigated the effects of loading rate on discont im; - ous deformation behavior and tensile properties.

In Advances in Cryogenic Engineering - Materials

,

vol. 34, Plenum.

New York, 1988, pp . 233-240.

295
EXPERIMENTAL PROCEDURES

Materials tested in this study were commercial AISI 304L, 310, and

316LN, typical austenitic stainless steels for low- temperature applications. Table 1 shows the chemical compositions of these alloys. The

materials were in the as -received, annealed condition. The AISI 316LN

was solution-heat- treated at 1423 K for 2 h. The grain sizes of AISI

304L, 310, and 316LN were 70, 140, and 100 ^m, respectively.

Table 1. Chemical Composition of the Steels Tested in This Study (wt.%)

C Mn P S Si Cr Ni Cu Mo N Fe

AISI 304L 0.02 1.4 0.02 0.01 0.6 18.4 9.7 — — — bal

.

AISI 310 0.08 1.7 0.02 0.02 0.7 24.8 20.8 0.1 - - bal

.

AISI 316LN 0.02 1.8 0.02 0.01 0.3 16.8 13.9 — 2.2 0.099 bal.

The round specimens had gauge lengths of 38 mm and reduced- section

diameters of 6.35 mm. Specimen strain was measured with two pairs of

clip- on strain gage extensometers; each pair included one normal and one high- range extensometer . The normal extensometer , with a 2. 5 -mm span and

a gauge length of 25.4 mm, was used to measure yield strength and initial

discontinuous deformation strain of the specimen. The high-range extensometer, with a span of 17.5 mm and a gauge length of 38 mm, was used to

measure specimen strain throughout the tensile test. Figure 1 shows a specimen with normal and high-range extensometers.

We used a programmable servohydraulic testing machine (having a maximum load of 250 kN) with a cryostat for the liquid helium tests.

With a high-resolution function generator, we varied loading rates from

0.5 to 5000 N/s

.

The machine was programmed for a ramp function in a load-control mode; that is, the machine uniformly increased the load on

the specimen until failure. The strain rate in the linear-elastic region

of the stress—strain curves was calculated to be 9 x 10" 8 s' 1 to 9 x

10" 4 s" 1

, corresponding to the loading rates of 0.5 N/s to 5000 N/s. To

shorten the testing time in 5 -N/s and 0.5 -N/s tests, we adopted a loading

rate of 50 N/s up to 80% of yield strength and stopped the test after one discontinuous deformation. We determined the dither ratio to be less

than 0.1% of the load signal, and from the oil pressure and oil flow, we estimated the actuator velocity of the testing machine to be 50 mm/s

.

We also performed displacement-control tensile tests with the same testing machine in a stroke -control mode and with the same extensometers

to compare the data obtained from load control and displacement control.

The strain rate of this tensile test was 2.2 x 10" 4 s' 1

. All these tensile tests were carried out with specimens submerged in liquid helium.

296
Figure 1. Specimen with normal and high- range exteosoaeters .

Load and strain were recorded on an X—Y recorder and a high-speed

digital memory recorder (sampling rate of 500 ^s) . During load- control

testing, we measured strain rate just before discontinuous deformation.

Using a bar-magnet, torsion-balance commercial magnetometer, we calculated the amount of a' -martensite in the deformed specimens.

RESULTS

Load—displacement curves for AISI 310 obtained from both load- and

displacement-control tensile tests are shown in Figure 2. In the figure

0.5 mm/min represents displacement-control test; the broken line indicates that the curve was calculated from a stroke signal because spec Iter,

extension exceeded the range of the high- range extetscmeters . No serrations"*’ were observed in load-control tests. At the lowest loading rate

0.5 N/s, the specimen failed at the first discontinuous deformation

At the loading rate of 5 N/s, the specimen deformed about 15 mm. in strain, within one deformation, and then failed. At 50 Ns, the

xIn the displacement-control tests, discontinuous deformation cons Lit;

of load drops, which we call serrations to distinguish them from the

discontinuous deformation the load-control tests.

297
Figure 2. Load—displacement curves obtained at

various loading rates for AISI 310.

deformation stopped, and the load increased again. With the increase of

loading rate, the stress to initiate discontinuous deformation (initiation strength) decreased. At 500 N/s , the deformation occurred just

above the yield strength. The number of deformations in the curve increased. At the highest loading rate, 5000 N/s, no apparent discontinuous deformation was observed owing to specimen heating.

Figure 3 shows the load—displacement curves for AISI 316LN and 304L

steel. The curves for AISI 316LN are similar to the curves for AISI 310.

The curves for 304L are different from those of AISI 310: the deformation stopped frequently, even at lower loading rates.

During load-control tests, the strain rate increased as the load

increased. Table 2 presents the strain rate just before the first discontinuous deformation.

Since the load was kept almost constant during the discontinuous

deformation in load-control tests, we could measure the exact velocity

of the deformation. In Figure 4, the monitoring screen shows the loadvs. -strain change obtained from AISI 310 at 5 N/s. Curve 1 is the load

and curve 2 is the strain change at the first (and the last) deformation.

The amount of this deformation was a 15 -mm change in length, and it lasted

0.575 s. The average velocity was 26 mm/s (6 x 10* 1 s" 1 in strain rate);

the maximum velocity was 32.5 mm/s (7.5 x 10* 1 s' 1 in strain rate), a value close to the maximum actuator velocity. So, at the higher deformation velocity, the load decreased during the deformation.

298
50

Figure 3. Load—displacement curves obtained at various loading rates

for (a) AISI 316LN and (b) for AISI 304L.

Table 2. The Strain Rate Just before the First Discontinuous Deformaticr

Loading Rate

(N/s)

Strain Rate (10‘ 4 s*

AISI 310 AISI 316LN AISI 304L

0.5 0.13 0.14 0.1

5 1.6 1.0 0.7

50 CM CM 3.5 1 500 5.7 13 6

299
The mechanical properties of AISI 310, 316LN, and 304L are shown

in Figures 5, 6, and 7 as a function of loading rate. In the figure,

S u , S^, Sy, El, and RA represent ultimate tensile strength, initiation

strength, yield strength, elongation, and reduction of area, respectively. Ultimate tensile strength and initiation strength decreased with

Figure 4. Load and strain during the deformation

of AISI 310 at 5 N/s

.

Figure 5. Tensile properties as a function

of loading rate for AISI 310.

300
an increase in loading rate. At the loading rate 5000 N/s , the ultimate

strengths of AISI 310, 316LN, and 304L were 71, 66, and 65%, respectively,

of those obtained from displacement-control tests. Yield strength was not affected by the loading rate. In load-control tests, elongation

decreased for AISI 316LN and most AISI 310 specimens, but it was not

significantly different for AISI 304L.

The amounts of a' -martensite in deformed specimens are listed in

Table 3. No a ' -martensite was detected in the threaded section of any

specimen and or in any part of the AISI 310 specimens.

Figure 6-. Tensile properties as a function of loading rate

for AISI 316LN.

Figure 7

.

Tensile properties as a function of loading rate

for AISI 304L.

Table 3. Amounts (vol.%) of a

' -martensite in Deformed Specimens

Location* AISI 304L

0 . 5mm/min 5 N/s 50 N/s

AISI 316LN

0.5mm/min 5 N/s 50 N/s

Section A 100 73 88 47 21 5

Section B 90 78 79 48 3 5 Section C 99 73 56 40 3 2

* A: close to the fracture surface

B: 15 mm apart from the fracture surface

C: close to the radius of the specimen reduced section

301
DISCUSSION

Deformation Behavior

We observed discontinuous deformation except at the highest loading

rate. At 500 N/s for AISI 310, the deformation occurred just 3% above

the yield strength, and the amount of the deformation reached about

0.7 mm, about 2% in strain. The strain rate just before the deformation

was 5.7 x 10* 4 s* 1

, which is usual in displacement-control tensile tests.

The serrations that occurred near yield strength seemed to be negligible.

A large amount of discontinuous deformation is characteristic of loadcontrol tests, as Ogata et al. reported in previous papers. 1 ’ 2 The initiation strength increased owing to less specimen heating at the lower

strain rate. The change in length during the first deformation for AISI

310 at loading rates of 500, 50, and 5 N/s was about 2, 7, and 15 mm. The deformation locally heats the specimen and stops when a balance oc- curs between the local strength and the work-hardening at the high temperature. Subsequent discontinuous deformations occur in other areas of

the specimen where less local strain hardening has occurred.

During the large deformation (about 15 mm) of AISI 310 at 5 N/s,

the entire specimen length continued to deform accompanied by continuous heating during the deformation. During this deformation, the work

done was about 30 kN x 15 mm - 450 J, and the power was 450 J/0.575 s -

783 J/s. The normal cooling capacity of this kind of specimen is about

3 J/s, less than the heat generated. The maximum temperature rise was estimated and measured as more than 100 K. 2

In displacement-control tests, the deformation is suppressed by the

testing machine, and when the deforming stress decreases, this suppression results in the serrations of smaller strain increments.

In AISI 304L, the deformation stopped frequently. This was due to

its lower initiation strength and also to its high work-hardening accom- panied by martensitic transformation. During the deformation, unstable

deformation continued because of the transformation.

The temperature rise in the specimen also determines the amount of

a' -martensite formed, which was less in the load-control tests than in

the displacement-control tests. The effect was significant in the load- .. control test of AISI 316LN at 50 N/s, though its total elongation was 60%

of that in the displacement-control test. We also observed that the

specimen deformed less uniformly in the load-control test owing to the

temperature gradient at large strains, where the temperature is the

highest

.

Tensile Properties

Yield strength did not change with the loading rate because the

strain rate in the elastic range was lower than 10' 3 s

=1

, and specimen

heating occurred after the plastic deformation began.

302
In load-control tests, ultimate tensile strength decreased 65 to

71% at 5000 N/s and 68 to 80% of that in the usual displacement- control

tests. This was due to the discontinuous deformation that leads to

specimen failure at higher temperature before the specimen is fully work

hardened at 4 K. This decrease was larger than that obtained in higher

strain- rate displacement-control tests.

3

This phenomenon indicates that

if a material at low temperatures is in the freely deforming condition,

it will fail at a much lower stress than the stress obtained in the usual

displacement -control tensile tests.

The strength to initiate discontinuous deformation decreases with

increases in the loading rate; this is closely related to the effect of

strain rate on the serrations in displacement-controlled tests. 8 The

initiation strength in displacement-control tests corresponds to the

range 50 to 500 N/s in load-control tests. The initiation strength in

load-control tests for AISI 310 and 316LN equals the ultimate tensile

strength at 0.5 and 5 N/s.

In AISI 304L, no effect of either the loading rate or control mode

on elongation was observed. This was due to lower initiation strength

and higher work-hardening rate supported by the a' -martensite transformation of the alloy.

For both 304L and 316LN alloys, the elongation to fracture was essentially the same for the 5 and 50 N/s load rates (Figures 6 and 7)

.

Therefore, with increasing load rate, the amount of a' -martensite per

unit strain increased in 304L and decreased in 316LN. This is similar to the trend observed by Reed and Walsh9 in these same alloys in

displacement-control tests at 4 K. They attributed the trend to specimen heating. Heating produced more a' -martensite in 304L because more

martensite forms in this alloy at higher temperatures; conversely, in

alloy 316LN there is no C-curve behavior and less a' -martensite forms

at higher temperatures. The explanation for the increased a' -martensite in the 316LN during displacement -control tests is the much larger

elongation.

CONCLUSIONS

1. Yield -strength remained constant with loading rate and reduction of

area decreased slightly.

2. Ultimate tensile strength decreased with increased loading rate,

reflecting specimen heating at a higher rate, to 65% of the strength

obtained in displacement-control tests.

3. The stress at which the first discontinuous deformation occurred with

increasing load rate was equivalent to the yield stress at higher

rates. The amount of deformation that occurred just above yield

strength reached 2%.

303
4. Large deformation (15 mm) occurred in AISI 310 at 5 N/s . During

this deformation, the entire specimen length continued very rapid

deformation with specimen heating within one load drop.

5. Average deformation velocities of 26 mm/s (7.5 x 10" 1 s" 1

) were

measured.

6. Measurement of a-martensite in the specimens after fracture

indicated temperature increases during deformation.

ACKNOWLEDGMENTS

This study was partly sponsored by Office of Fusion Energy, U.S.

Department of Energy.

REFERENCES

1. T. Ogata and K. Ishikawa, Cryogenics 26:365—369 (1986).

2. T. Ogata, K. Ishikawa, 0. Umezawa, and T„ Yuri, Discontinuous

deformation during tensile tests in liquid helium, submitted to

Cryogenics . 3. T. Ogata, K. Ishikawa, and K. Nagai, Tetsu- to-Hagane 71:1390—1397

(1985), in Japanese.

4. D. T. Read and R. P. Reed, Adv. Crvog. Eng . 26:91 (1980).

5. A. Seeger, Dislocation and Mechanical Properties of Crystals . John Wiley, New York (1957), p. 206

6. Z. S. Basinski, Proc. R. Soc . A240:229 (1957).

7. G. Y. Chin, W. F. Hosford, Jr., and W. A. Backofen, Trans . AIME

230:1043 (1964).

8. R. P. Reed and R. P. Walsh, National Bureau of Standards, Boulder,

Colorado, unpublished data.

9. R. P. Reed and R. P. Walsh, Tensile strain rate effects in liquid

helium Adv. Crvog. Eng. 34:199 (1988).

304
ROUND -ROBIN TENSILE AND FRACTURE TEST RESULTS

FOR AN Fe—2 2Mn—1 3 Cr—5Ni AUSTENITIC STAINLESS STEEL AT 4 K*

H. Nakajima, K. Yoshida, and S. Shimamoto

Japan Atomic Energy Research Institute

Tokai-mura, Ibaraki-ken, Japan

R. L. Tobler, P. T. Purtscher, and R. P. Reed

National Bureau of Standards

Boulder, Colorado

Round-robin tests are a part of the U.S.—Japan cooperative

program to develop standard tensile and fracture test procedures for austenitic stainless steels at 4 K. The first round

of tests for an Fe—22Mn—13Cr—5Ni steel are described in this

paper. The interlaboratory scatter in 4-K measurements is re- ported and relevant aspects of test procedure are discussed.

INTRODUCTION

Cryogenic standards are needed to support magnetic fusion- energy

developments. At the United States—Japan meeting on Low Temperature

Structural Materials and Standards in December 1984, the development

of property measurement standards was cited as a first step towards the

ultimate goal of establishing design standards for fusion-energy structures. A program was therefore initiated to evaluate and standardize

the existing laboratory procedures. This paper describes the first round

of interlaboratory tests that are now under way in the United States and

Japan.

APPROACH

Seven laboratories participated, five from Japan and two from the

United States. The participants are volunteers from government, aca- demic, and industrial organizations with recognized materials testing

experience. Table 1 identifies the laboratories, their abbreviations,

and associated personnel.

In Advances in Cryogenic Engineering — Materials , vol. 34, Plenum,

New York, pp . 241-249.

305
Table 1. Participating laboratories and personnel

Institute Supervisor Researcher Staff

Hitachi Research

Laboratory (HRL)

T. Matsumoto Y . Wadayama H. Sato

Japan Atomic Energy

Research Institute (JAERI)

S. Shimamoto H . Nakaj ima

K. Yoshida

M. Oshikiri

Kobe Steel, Ltd. (KSL) T. Horiuchi M. Shimada

S . Tone —-

Lawrence Livermore

National Laboratory (LLNL)

and Lawrence Berkeley

Laboratory (LBL)

J. Miller and

J.W. Morris, Jr.

L.T. Summers J.W. Chan

M. Strum

National Bureau of

Standards (NBS)

R.P. Reed R.L. Tobler

P.T. Purtscher

R.P. Walsh

National Research Institute for Metals (NRIM)

K. Ishikawa T. Ogata

K. Nagai

T. Yuri

Tohoku University,

Reasearch Institute for

Strength and Fracture

of Materials (RISFM)

H . Takahashi T . Shoj

i

K . Nakano

At the time of our study, standard 4-K tensile and fracture test

procedures did not exist. Each laboratory used its own procedures. In

general, they followed the established room- temperature standards, 1 ' 8 except where modifications were necessary for tests at 4 K.

The test material was a newly developed Fe—22Mn—13Cr—5Ni austenitic

stainless steel; its composition, in wt.%, is Fe—21 . 79Mn—12 . 82Cr—4 . 94Ni—

0 . 05C—0 . 36Si—0 . 013P—0 . Q05S—0 . 212N. It was supplied as 70-mm- thick plate

from a 15 -metric -ton industrial heat, in the annealed condition. Samples

were cut from locations at the midthickness of the original plate. One

laboratory (JAERI) machined all the test specimens to minimize variations

in preparation. The number of specimens distributed can be inferred from

the results in the text. A questionnaire was distributed with the speci- mens to identify the individual laboratory apparatus and procedures

.

306
A

Thickness, B = 25.4 mm

Figure 1. Specimen configurations used in this study: A) tensile specimen, and B) compact specimen. (All dimensions are in mm.)

SPECIMEN CONFIGURATIONS AND PREPARATIONS

Tensile Tests

Cylindrical tensile specimens machined with the load axis parallel

to the long- transverse plate direction were used by everyone, but the

specimen dimensions were varied to suit individual laboratory preferences. Figure 1 defines the relevant specimen parameters, and Table 2 gives the specimen dimensions for each laboratory. The conventional

gage- length- to -diameter (G/D) ratios are 4 or 5 for cylindrical specimens. 2-5 In this study, three laboratories used G/D - 4 and marked the

gage length on the reduced specimen length by means of punch or scribe

marks. Other laboratories did not mark a gage length but calculated

elongation from changes in overall specimen length or distance between

fillets before and after testing (see Table 2)

.

Fracture -Toughness Tests

A compact specimen was used in this study, as shown in Fig. 1. All

specimens were identical in thickness (B = 25.4 mm), width (W - 50.8 mm),

and notch length (an = 27 mm). The notch orientation was TL, and displacement was measured at the loadline.

The specimens tested in Japan were fatigue -cracked by one laboratory

(KSL) in air, at 295 K, and at a cycle frequency of 40 Hz. The desired

initial crack length, a0 , was 32 mm. The crack was first extended 1.7 mm from the notch by using maximum and minimum fatigue loads of 17.6 and

0.78 kN; further crack extension from 1.7 to 3.5 mm was then accomplished

by using loads of 9.8 and 0.78 kN. The specimens tested in the United

307
Table 2. Tensile specimen dimensions (see Fig. 1 for definition of

symbols)

Institute D

(mm)

G

(mm)

A

(mm)

G/D Interval Used for

Elongation Measurements

HRL 7.00 28.0 40.0 4 gage marks , punched

JAERI 7.00 - 42.4 _ total specimen length

KSL 7.00 - 42.4 - total specimen length

LLNL/LBL 6.35 25.4 33.5 4 extensometer screw indentations

NBS 6.25 25.4 38.1 4 gage marks , inked

NRIM 6.25 - 32.0 - total specimen length

RISFM 6.35 — 31.75 — distance between ends of fillets

States were fatigue -cracked in liquid nitrogen at 76 or 77 K; the load

ratio was 0.1; the maximum stress- intensity factors were less than

30 MPa-mh. After precracking, NBS side -grooved their specimens to a net

thickness reduction of 20%. The side grooves were 3.8 mm deep with an included angle of 60° and a notch- tip radius of 0.4 mm. The different

precracking temperatures used in the two countries had no measurable

effect on the Jj c measurements.

TEST PROCEDURES

Tensile Properties

The tensile properties of interest are the yield strength (YS) , the

ultimate strength (UTS) , the reduction of area (RA) , the elongation (EL)

,

and Young's modulus (E) . Procedures for testing at 295 K are specified

in JIS and ASTM standards

.

1 " 8 Load—extension or load—time curves or both

are recorded. Yield strength is measured at 0.2% plastic strain by an offset method applied to load- strain- gage , load- extensome ter , or timechart recordings. Ultimate strength is obtained by dividing the maximum

load by the original area. Elongation is computed by dividing the change

in length by the original length. Reduction of area is obtained by dividing the change in cross-section area by the original area. The slope

of the stress—strain curve in the elastic region gives Young's modulus.

Most of the tests were conducted using screw-driven test machines with

crosshead displacement control at various speeds. Table 3 lists the

nominal strain rates calculated by dividing the free -running crosshead

speed by the initial length of the reduced section.

308
Table 3. Summary of tensile test techniques

Institute Crosshead Speed

(mm/min)

Strain Rate

( 10" 4 s" 1 ) Measurement Method for

0 . 2%-0ffset YS

HRL 0.5, 1.0a 2.1, 4.

2

a strain gage

JAERI 0.6, 3 .

0

a 2.4, 11.

8

a strain gage

KSL 12b

3.9

7 .

8

b

strain gage

LLNL/LBL 0.35 1.7 extensometer

NBS 0.2, 0 .

5

a 0.9, 2 .

2

a strain gage , time

,

extensometer

NRIM 0.5 2.6 extensometer

RISFM 1 5.2 strain gage, time

Notes: a.

b.

Rate change was used: the first number is the rate during YS

measurement; the second, the rate during UTS measurement.

Different speeds were used in the second test.

Fracture Properties

The fracture properties of interest are Jjc and Kj c (J), estimated

from the expression Kj c

(J) - (Jjc’E)^. The test procedure of interest is

the computer-aided single -specimen unloading compliance technique. 8 During the test, the specimen is unloaded periodically, and its stiffness is

measured from the slopes of the unloading lines. From changes in stiffness, the crack extensions, Aa, are inferred and plotted versus J, and

Jj c is obtained using Method E 813-81 guidelines. 7 That standard imposes

no controls on hardware or software. In this study, most tests were con- ducted using servohydraulic machines in stroke control, but two laboratories used servohydraulic machines in clip-gage-displacement control.

Also, one laboratory used specimen rotation and crack growth correction

factors to their J—Aa calculations, but other laboratories didn't. The

computed initial and end-of-test crack lengths, a 0 and ap , were compared

with direct optical measurements after testing.

309
RESULTS

Tensile Properties

Figure 2 shows a typical 4-K tensile-test record. Serrations oc- curred in the plastic -strain range, which is usual for ductile alloys at

4 K. In this steel, the serrations begin at plastic strains near 6% and

do not interfere with YS determinations at 0.2% strain.

The YS and UTS values, plotted versus nominal strain rate in Fig. 3,

indicate no measurable effect for the range of strain rates covered in

this study, 0.9 to 7.9 x 10" 4 s' 1 for YS and 2.1 to 11.8 x 10' 4 s' 1 for

UTS. Likewise, no size effect is evident. Therefore, the measurements

for all fourteen tests were pooled; a summary is given in Table 4. The

means and ranges are

YS - 1242 MPa; + 5.0%, - 5.3%

UTS - 1667 MPa; + 2.7%, - 4.2%

EL - 36.0; + 6.0%, -5.0%

RA - 49.0; + 5.0%, -5.0%

E - 200 GPa; + 7.5%, - 6.0%.

Figure 2. Time-chart record for the Fe—22Mn—13Cr—5Ni steel at 4 K.

310
1800

<o CL

1700

1600 -

1500

cocn® 1400 -

CO

1300 -

1200

1100

o

o

A

A

A

6

A 8

YS TS D A

A A HRL (7.00 X 40.0) • O JAERI (7.00 X 42.4) A A KSL (7.00 X 42.4) O LLNL / LBL (6.35 x 33.5) V NBS (6.25 X 38.1) D m NRIM (6.25 X 32.0) RISFM (6.35 x 31.75)

llllll I 1 I I 1 I 111 1 1 I II I 111 | L_L_L

10

-4 10 -3 10

-2

Strain Rate (s ^

Figure 3. Yield and ultimate strength measurements at 4 K.

The YS values in Table 4 represent strain- gage or (in two cases)

extensometer techniques. Time-chart data were also reported by some laboratories. When the 0.2% -offset procedure was applied to the time

charts, the YS values were nearly equivalent or slightly higher than

values from strain or extensometer plots, with differences ranging from

—2 to 6%.

Fracture Toughness

Fourteen Jjc tests were attempted. One specimen (A17) failed owing

to computer malfunction and another (A19) was evaluated using a nonstandard method, 9 which will be considered in a separate paper. A typical

load—displacement record and resistance curve for a stroke-control com- pliance test is shown in Fig. 4, and the Jj c data are listed in Table 5. A total of twelve tests were performed by seven laboratories using the

compliance technique. The measurements for specimens precracked at 76 K

and side-grooved are indistinguishable from the bulk of results for

specimens precracked at 295 or 77 K and not side -grooved. Two specimens

(A3 and A4) , however, gave low toughness values, which could be due to

material variability. A summary, including these values, is given in

Table 5. The mean values and ranges are

J Ic - 135 kJ/m2

; +23%, - 35%

and KIc (J) = 166 MPa-m4; +8%, - 19%.

311
Table 4. Tensile test results

(YS from strain gage,

for Fe—

:

, except

22Mn—13Cr—5Ni

as noted)

steel at 4 K

Insti- Spec

.

YS UTS EL RA E

tute (MPa) (MPa) (%) (%) (GPa)

HRL D1 1238 NA NA NA 192

D2 1176 1597 36 48 193

JAERI B1 1304 1686 31 50 208

B2 1260 1688 33 51 204

B5 1227 1633 37 50 205

B6 1227 1643 37 49 205

KSL B3 1248 1656 35 46 215

B4 1261 1648 36 53 204

LLNL/LBL A2 1181 a 1613 36b 29b 194

A12 1254a 1651 32b 26b 195

NBS A18 1281 1713 37 50 201

A19 1275 1676 30b 48b 205

NRIM Cl 1230a 1700 35 47 188

C2 1230a 1690 42 44 192

RISFM El 1202 1697 49 d 51 196

E2 1276° 1710 39 54 NA

SUMMARY:

Number of tests 16 15 11 12 15

Maximum value 1304 1713 42 54 215

Minimum value 1176 1597 31 44 188

Arithmetic mean 1242 1667 36 49 200

Standard deviation 35.5 35.5 2.9 2.8 7.3

Notes: a. extensometer value

b. broke on gage mark, not used in average

c

.

time -chart value

d. not used in average

312
Figure 4. A load-vs-displacement test record and the resulting resistance

curve for a Jj c test at 4 K, showing results obtained using

servohydraulic machines operated in stroke control.

Table 5. Fracture test results for Fe—22Mn—13Cr—5Ni steel at 4 K

Institute Spec

.

J Ic

(kJ/m2

) KlcW ao (MPa -in*)

Agreement

(%) a

p Agreement

(%)

E

(GPa)

HRL A21 166 179 -0.9 -8.7 193

JAERI A13 138 169 -1.5 -30.1 208

A14 147 174 -1.8 -17.8 208

KSL A15 146 178 0.3 NA 215

A16 143 175 -4.3 NA 215

LLNL/LBL A4 103 a 148 a -2.2 -14.3 194

A5 140 172 -1.8 -16.7 194

A6 153 180 -2.8 -10.3 194

NBS Al 127 162 1.5 -16.7 201

A2 140 170 1.0 -8.3 201

A3 88 a 135 a -0.7 —4 .

8

201

NRIM A17 (computer malfunction)

A18 125 154 -3.0 -10.0 190

RISFM A19 144b 168b NA NA 196

SUMMARY:

Number of tests 12 12

Maximum value 166 180

Minimum value 88 135

Arithmetic mean 135 166

Standard dev. 21.5 14

Notes: a. Low value, possibly due to material variability (see text)

b. Key curve data, not included in the average (see text)

313
DISCUSSION

This interlaboratory testing of structural alloys is the first

ever conducted at 4 K. The measurements show fair agreement despite the

lack of formally written 4-K test procedures. Further refinements are possible

.

Tensile Tests

Scatter in tensile measurements was highest for E (± 7%) and lowest

for UTS (± 4%). In general, the scatter was slightly higher than that

reported for a room- temperature study. 10 The scatter in EL may have been

increased by the variety of gage -length marking techniques used. Gagemarking techniques should be improved, since three out of six specimens

failed at gage marks when gage marks were made on the reduced section.

Concerning YS measurements, this study indicates that perhaps a small bias may be introduced (possibly dependent on the test apparatus)

when the 0.2% -offset procedure is applied to time -chart plots, rather

than extensometer or strain- gage plots. Time charts usually show some nonlinearity at the start of testing, as well as other irregularities

(Fig. 2). For convenience, however, the time-chart method is quite useful in alloy screening or commercial production control, where high precision is not demanded.

Fracture Tests

The fracture- toughness measurements showed fair agreement. The

scatter of Jjc for 12 compliance tests (up to 35%) was greater than the

± 17% scatter in a room- temperature study. 11 If our two lowest values

are eliminated, with the assumption that they reflect material variability, then the mean values are increased and the ranges are reduced

considerably to

J Ic - 142.5 kJ/m2

; +.16.5%, - 12%

and KIc (J) - 171 MPa-m^; +5%, - 10%.

The compliance procedure therefore appears to be a reliable method for

Jjc determinations at 4 K. Some problems that surfaced in this study

should be addressed before standardization. These include the negative

crack- growth effect, the disagreement between measured and computed endof-test crack lengths, and the choice of E modulus.

As shown in Fig. 4A, the resistance curve at the onset of testing

shows negative Aa (i.e., an anomalous decrease of crack length). At

this stage of development, we treat this by shifting the blunting line

to the left on the Aa axis until it coincides with the upper portion of

the J—Aa trend, but this procedure is arbitrary and not recognized in

Method E 813-81.

314
As shown in Table 5, most tests for specimens without side grooves

did not give good ap predictions; the measured and calculated values

must agree to within 15% according to Method E 813-81, section 9.3.4,

for room- temperature testing. The NBS measurements for side -grooved

specimens meet this requirement, as suggested in Methods E 813-81,

section Al .4.

Young's modulus must be measured or estimated before Jj c tests. Any

uncertainty affects the resistance curve, because Aa is inferred from the

normalized compliance, EB5/P. The conversion from Jjc to Kj c (J) is also

affected. This was assessed at JAERI in a post-test analysis by substituting various values of E and recalculating J—Aa, Jjc , and Kj c (J). We

found that Kjc

(J) increases 10 MPa-m^ for a 10 GPa increase of E. In

Table 4, the reported moduli range from 188 to 215 GPa. Thus, the spread

of Kj c (J) data would have been reduced by about ± 7% if everyone had used

the same value of E. Since Young's modulus measurements at 4 K are cov- ered by Method E 111, we recommend that future interlaboratory tests

incorporate that method and specify one Young's modulus for J-integral

tests

.

CONCLUSIONS

Seven laboratories performed smooth-bar tensile and Jj c tests of an austenitic stainless steel at 4 K. The tensile results show good agreement, independent of minor variations in specimen size and strain rates

below a value of 10‘ 4 s" 1

. The tensile property measurement ranges were

± 7% or less and not much greater than would be expected for tests at

295 K. Also, the unloading- compliance test appears to be a reliable

method for Jjc measurement at 4 K. The range of Jjc measurements was within 35%, or 16.5% if the two lowest values are discounted.

ACKNOWLEDGMENTS

The organizers of this activity wish to express sincere appreciation

to all the participants for their contributions. Without such collaboration, the development of consensus standards is impossible.

The study was supported, in part, by the Office of Fusion Energy,

U.S. Department of Energy.

315
REFERENCES

1. Test Pieces for Tensile Test for Metallic Materials, Japanese

Industrial Standard, Designation JIS Z 2209-1980. English translation, Japanese Standards Association, Tokyo (1980).

2. Method of Tensile Test for Metallic Materials, Japanese Industrial

Standard, Designation JIS Z 2241. English translation, Japanese

Standards Association, Tokyo (1980).

3. Standard Methods and Definitions for Mechanical Testing of Steel

Products, Designation ANSI/ASTM A 370-77, 1986 Annual Book of ASTM

Standards, vol. 03.01, American Society for Testing and Materials,

Philadelphia (1986)

.

4. Standard Methods of Tension Testing of Metallic Materials, Designation E 8 -85b, 1986 Annual Book of ASTM Standards, vol. 03.01,

American Society for Testing and Materials, Philadelphia (1986).

5. Standard Methods of Tension Testing of Metallic Materials (Metric),

Designation E 8M-85, 1986 Annual Book of ASTM Standards, vol. 03.01,

American Society for Testing and Materials, Philadelphia, (1986).

6. Standard Test Method for Young's Modulus, Tangent Modulus, and Chord

Modulus, Designation E 111-82, 1986 Annual Book of ASTM Standards,

vol. 03.01, American Society for Testing and Materials,, Philadelphia (1986).

7. Standard Test Method for Jjc A Measure of Fracture Toughness, Designation E 813-81, 1986 Annual Book of ASTM Standards, vol. 03.01,

American Society for Testing and Materials,, Philadelphia (1986).

8. G. A. Clarke, Single -Specimen Tests For Jjc Determination —

Revisited, in: Fracture Mechanics: Thirteenth Conference, ASTM STP

743, R. Roberts, Ed., American Society for Testing and Materials,

Philadelphia (1981)

.

9. M. Takeuchi, T. Shoji, H. Takahashi, and T. Anayama, Reports of the

Research Institute for Strength and Fracture of Materials, Tohoku

University, Sendai, Japan, 17, 1-12 (1984).

10. R. H. Heyer, ASTM Bulletin, Technical Publication TP 83, American

Society for Testing and Materials,, Philadelphia, 57-62 (1949).

11. G. A. Clarke, J. D. Landes, and J. A. Begley, Journal of Testing and

Evaluation, 8, 221-232 (1980).

316
SIZE, SIDE -GROOVING, AND FATIGUE -PRECRACKING EFFECTS

ON J- INTEGRAL TEST RESULTS FOR SUS 304 STAINLESS STEEL AT 4 K*

M . Shimada

Kobe Steel

Kobe, Japan

R. L. Tobler

National Bureau of Standards

Boulder, Colorado

T. Shoji and H. Takahashi

Research Institute for Strength

and Fracture of Materials

Tohoku University

Sendai, Japan

Tests were conducted to investigate specimen preparation

effects on the J- integral fracture- test procedure at 4 K.

Proportional compact specimens of a fully plastic steel were

tested using the compliance technique. The Jjc measurements

were not affected by side grooving, but decreased slightly as specimen thickness increased from 12.5 to 25.4 mm. The resistance curve slope decreased significantly with both thickness

and side grooving. Fatigue -precracking and sustained- loading

effects are discussed.

INTRODUCTION

The 4-K fracture toughness of austenitic alloys is usually determined by J- integral tests using compact specimens and the unloadingcompliance technique. The compliance test procedure was developed for

ferritic steels at room temperature, 1-3 and we applied it later 4

' 7 to

austenitic stainless steels at 4 K. Because of dissimilar material

behavior (i.e., phase transformations, serrations, and adiabatic heating), several aspects of the 4-K procedure warrant verification. This

In Advances in Cryogenic Engineering — Materials , vol. 34, Plenum,

New York, 1988, pp

.

251-258.

317
paper investigates the influence of sustained loading, specimen size,

side -grooving, and precracking temperature on results for a hightoughness SUS 304 steel. This work was part of the United States—Japan

Cooperative Program for 1986 and was conducted to contribute to the

development of a standard 4-K test procedure.

MATERIAL AND TEST PROCEDURES

The material selected for these tests was a research heat of type

SUS 304 austenitic stainless steel with 4-K yield strength and fracture

toughness values of 615 MPa and 300 MPa-m\ This steel exhibits martensitic phase transformations when deformed at 77 or 4 K, but not at 295 K,

and it lies at the high- toughness extreme of the trend for the annealed

AISI 304 alloys. 4 A 30-mm-thick plate of this steel was obtained and

tested in the annealed (1080°C) and water -quenched condition. The com- position and tensile properties are given in Tables 1 and 2.

Three specimen configurations were studied: standard 25-mm-thick

compact specimens (1TCT) , with and without side grooves, and proportionally reduced 12 . 5 -mm- thick compact specimens (HTCT) without side grooves.

Basic dimensions are shown in Fig. 1. The width- to- thickness ratios

(W/B) were 2.0, the notch orientations were TL, and razor blades for

clip-gage attachment were located on the loadline. 1

Table 1. Chemical composition for SUS 304 steel (wt.%).

Fe C Cr Ni Si Mn P S

bal

.

0.06 18.23 12.95 0.58 1.21 0.027 0.006

Table 2. Mechanical properties for SUS 304 steel.

Test Yield Ultimate Elonga- Red. of Young'

s

Brinell

Temperature Strength Strength tion Area Modulus Hardness

(K) (MPa) (MPa) (%) (%) (GPa)

295 215 529 58 — — 131

4 615 1475 57 52 217 —

318
62

Fig. 1. Compact specimens used in this study. Dimensions are in mm.

The specimens were fatigue -precracked at 295 or 77 K to crack length

ratios (a/W) near 0.6, using a load ratio of 0.1 and a cycle frequency of

20 Hz. The maximum fatigue loads (Pmax) are listed in Table 3. After

precracking, two of the 1TCT specimens were side -grooved 2 mm deep on each face. The groove root radius was 0.5 mm and the net thickness re- duction was 16%. For the side-grooved specimens, compliance was calculated using the effective thickness (Be ff) and J was calculated using the

net thickness (Bnet ) , per ASTM Methods E 813-81 sections 8.7.1 and 9.1.1.

The specimens were tested at 4 K using a turret-disc computer-aided

unloading- compliance apparatus. 6 The load-vs . -displacement (P-5) curves

and resistance (J—Aa) curves were autographically recorded. Crackinitiation resistance, Jjc , is defined as the point of intersection of

the blunting and regression lines; crack-growth resistance, AJ/Aa, is the

resistance -curve slope. Since the apparent slope computed for a specimen

without side grooves may be inaccurate, 5 a corrected slope was obtained

from a line plotted between Jjc

* and the physically measured end-of-test

crack- extension increment, Aap. During tests, some specimens were held

for 10 min at fixed stroke positions near the limit load while the load

cell and clip-gage outputs were monitored for relaxation. After the un- loadings were completed, the specimens were heat- tinted and broken apart;

the Aap values were measured, and the fracture surfaces were examined by

scanning electron microscopy (SEM)

.

319
Table 3. Precracking conditions (295 and 77 K) and 4-K toughness results

Specimen Fatigue Precracking Fracture Toughness

Description Conditions

Type and Number Temp . Pmax p /p_ ^max,

(K) (kN)

rmax/ rL

( MPa-m4

) a/W J Ic KIC (J)

' (kJ/m2

) (MPa -in

5

)

HTCT, Al a 295 4.9 0 . 70d 34.2 0.603 343 e 272e plain A2 295 5.4 0 . 77 d 37.7 0.605 453 314

A3 77 4.9 0.33 35.1 0.610 434 307

A4 77 5.9 0.41 42.8 0.613 415 300

average - 434 307

1TCT

,

B4 295 13.7 0.48d 32.8 0.598 392 292

plain B5 295 10.8 0.37 25.7 0.595 373 285

B6 295 16.7 0 . 57d 39.4 0.594 392 292

B7b 77 13.7 0.22 34.1 0.606 —

—

B8 77 11.8 0.18 27.5 0.591 412 299

average - 392 290

1TCT

,

B1 295 11.8 0.42 29.3 0.605 386 289

side B2C 295 NA NA NA NA NA NA

grooved B3 77 13.7 0.19 29.5 0.571 397 294

average - 392 291

Notes

a — skewed crack front development

b — incomplete resistance curve

c — failure due to machine malfunction

d - final fatigue load cycle shows some hysteresis

e — not used in average

320
RESULTS AND DISCUSSION

Tables 3 and 4 display the principal results of this study. As we expected, the fracture behavior at 4 K was fully plastic. Representative

test records, resistance curves, and fractured specimens are shown in

Figs. 2, 3, and 4. Ductile fracture in this steel consisted of stretch

zone and dimple formation, as shown in Fig. 5.

Sustained-Loading Behavior

Drift in the load or displacement signals owing to creep can affect

compliance results by causing nonlinear unloading slopes. 1 This problem

is common for austenitic stainless steels at 295 and 77 K, where load re- laxations occur readily during the first few minutes of static holding in

J tests. In this study at 4 K, however, load relaxations did not occur. The SUS 304 steel of this study had a rather low yield strength at 4 K,

and because time -dependent plasticity effects are even less likely with

higher strength at this temperature, time -dependent effects are not ex- pected for most austenitic stainless steels at 4 K.

Table 4. Summary of resistance curve features for SUS 304 steel at 4 K

Regression Line Points in Accuracy

Specimen Type

Slope

Computed (Corrected)

Exclusion

Interval3

of Aa^.

Prediction0

(kJ/m2 /mm) (%)

HTCT, plain Al 470 (NA) 7 NA

A2 360 (170) 9 90

A3 500 (367) 9 57

A4 260 (196) 11 33

average 398 (243)

1TCT, plain B4 140 (NA) NA NA

B5 275 (254) 3 43

B6 230 (203) 5 10

B8 330 (297) 4 20

average 279 (251)

1TCT

,

B1 115 (120) 3 -5

side -grooved B2 NA NA NA NA

B3 195 (199) 2 -2

average 155 (159)

Notes : a — at least four points are required by E 813-81, section 9 .

3

b — agreement within 15% is required by E 813-81, section Al .

4

321
Displacement, d, mm

Fig. 2. Load—displacement records for specimens without side grooves:

A) 4TCT and B) 1TCT. Load is in metric tons.

322
J-Integral,

kJ

/

m

Crack Extension, Aa, mm

Fig. 3. J-Aa curves for SUS 304 steel at 4 K: A) HTCT, plain,

B) 1TCT plain, C) 1TCT side grooved, and D) comparison.

323
Fig. 4. Fractured 1TCT specimens, showing the heat-tinted crack

extension increments: A) without the side grooves,

crack tunneling; B) with side grooves, uniform growth.

Fig. 5. Fractograph (SEM) ; specimen B1 at 4 K.

324
Specimen Size

Specimens smaller than the routine 1TCT size are of practical interest for use in multiple-specimen cryostats or for characterizing stock

that has limited dimensions. For specimens precracked at 77 K, the mea- sured Jj c values for HTCT specimens were slightly higher than those for

the 1TCT size, as shown in Table 3. This difference was less than 10%

and, therefore, of minor significance. In contrast, the resistance curve

slope was more strongly affected, as shown in Table 4 and Fig. 3. The

HTCT specimens curves had the highest slopes, the most data points in the

exclusion interval, and the poorest agreement between the measured and

computed Aap values. Likewise, the resistance -curve slopes for plain

1TCT specimens were higher than those for side-grooved specimens; this

is partly due to crack tunneling, as discussed below. Another possible

explanation for size effects is a departure from J-controlled crack

growth. According to Method E 813-81, the specimen thickness and ligament (b) dimensions must satisfy certain minimum values to ensure Jcontrolled crack growth. The ligament sizes for our hTCT tests are borderline at about 9.9 mm, whereas 10 mm is the calculated requirement

(Method E 813-81, section 9. 3. 5. 3):

B,b > 25Jic/oy, (1)

where ay, the flow stress, is the average of the yield and ultimate

strengths

.

Side Grooves

The use of side grooves is a controversial option. Side grooves

increase machining costs but may be desirable to alleviate crack- front

tunneling in tough materials 1 or to achieve plane-strain conditions more

perfectly. In this study, side grooves on 1TCT specimens lowered the

apparent resistance -curve slopes to realistic values without affecting

Jj c

. There were several beneficial effects: the crack tunneling was reduced (Fig. 4) , the Aap was accurately predicted, and the regression

line slope was accurately computed. These results are not unexpected

(see Method E 813-81, section Al.4). What is new, however, is the observation that side grooves tended to enhance discontinuous deformation

at 4 K. Thus, load drops were more pronounced with side grooves than

without them, and side-grooves reduced the number of points within the

0.15- to 1.5 -mm exclusion interval. For this reason, our tests of the

side-grooved 1TCT specimens did not meet the requirement of obtaining at

least four valid J—Aa points within the interval (Method E 813-81, sec- tion 9.3). Since the serrations were less pronounced in smaller specimens, in retrospect it appears that tests using side-grooved specimens

of an intermediate thickness between the hT and 1TCT sizes would have

satisfied all the requirements of E 813-81.

325
Fatigue Precracking

Many investigators precrack their specimens at 295 K, assuming that

Jj c measurements at 4 K will be unaffected. No proof is available, however, and Method E 813-81 states no guidelines for precracking and fracturing at dissimilar temperatures. The rules (section 7.6 of Method E

813-81) state that Pmax should be less than 40% of the limit load, P^:

Pmax * 0.4 x PL - 0.4 x Bb 2 OY/(2W + a), (2)

or be such that

AK < E x 0.005, (3)

whichever is less. The above equation requires AK < 34 MPa-ni 2 for the

SUS 304 steel at 295, 77, and 4 K.

Differences in crack-tip plastic-zone size and microstructure do

exist for fatigue cracks formed at room or cryogenic temperatures. For

SUS 304 steel, the plastic zones at 295 K were fully austenitic, whereas

those at 77 or 4 K were partially martensitic. Also, the zones at 295 K

were larger. Assuming - 30 MPa'm^ and using Irwin's approximation,

R “ (Kmax/^y) (V6 *) (4)

we estimated the plastic zone radii (R) as 1.0, 0.20, and 0.12 mm at 295,

77, and 4 K, respectively. This shows that the differences in plastic

zones were clearly much greater between 295 and 77 K than between 77 and

4 K. In previous tests of AISI 310S, precracking at 76 or 4 K had no effect on resistance curves at 4 K. 5 Therefore, in this study we com- pared precracking at 295 or 77 K.

The following interpretation is offered for the results in Table 3.

For the hTCT specimens: the Jj c values for two specimens precracked at

295 K at high Pmax/RL were scattered; hence, the comparison with speci- mens precracked at 77 K is inconclusive. For the 1TCT specimens: the Jj c values for two specimens precracked at 77 K at low Pmax/PL ratios (0.18

and 0.19) were slightly higher than results for four specimens precracked

at 295 K at high Fjnax/RL ratios (0.37 to 0.57). This suggests • that precracking at 295 K may give lower Jj c values, but the effect (less than

5%) falls within measurement uncertainty. In these tests, we used high

precracking loads to accentuate any possible effects. For routine testing, lower load ratios (Fmax/RL

< 0*4-) anc* AK values (-25 MPa-mh) are normally used, in which case any effects of precracking at 295 K on Jj c measurements at 4 K should be negligible.

326
SUMMARY

Tests of an SUS 304 austenitic stainless steel were performed at 4 K

to evaluate the unloading-compliance Jj c test technique. The validity

criteria of the ASTM room- temperature standard were used to interpret the

data at 4 K. The following points are significant:

1. There is no evidence of load relaxation during static holding for

this material. This indicates that time -dependent plasticity effects

on the compliance test procedure are negligible at 4 K.

2. Jj c decreased 10% or less for a specimen- thickness increase from

12.5 to 25.4 mm, whereas the resistance -curve slope decreased rather

substantially with thickness, and with side grooving.

3. The Jj c measurements for 1TCT specimens were not affected by side

grooves. Side -grooving improved the accuracy of crack- growth predictions by the compliance method but reduced the number of valid

data points within the exclusion interval.

4. Limited data indicate that at fatigue precracking at 295 K does not

affect Jj c measurements at 4 K if the maximum fatigue loads are kept

within reasonable limits, as discussed in the text.

ACKNOWLEDGMENTS

This work was performed as part of the United States—Japan Cooperation Program in magnetic fusion energy for 1986. The research heat of

SUS 304 was supplied by Nippon Stainless Steel, and the fatigue precracking was performed at JAERI ; the authors greatly appreciate this

assistance. This study was supported, in part, by the Office of Fusion

Energy, U. S Department of Energy.

REFERENCES

1. Standard Test Method for Jj c , A Measure of Fracture Toughness, ASTM

Designation E 813-81, in: 1986 Annual Book of ASTM Standards, Section 3, vol. 03.01, American Society for Testing and Materials,

Philadelphia (1986).

2. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, SingleSpecimen Tests for Jj c Determination, in: Mechanics of Crack Growth,

ASTM STP 590, American Society for Testing and Materials, Philadelphia (1976).

3. G. A. Clarke, Single-Specimen Tests for Jj c Determination—Revisited

,

in: Fracture Mechanics: Thirteenth Conference, ASTM STP 743,

R. Roberts, ed. , American Society for Testing and Materials,

Philadelphia (1983).

327
4 . R. L. Tobler, D. T. Read, and R. P. Reed, Strength—Toughness Relationship for Interstitially Strengthened AISI 304 Stainless Steels at

4 K Temperature, in: Fracture Mechanics: Thirteenth Conference,

R. Roberts, ed. , ASTM STP 743, American Society for Testing and Materials, Philadelphia (1981).

5. R. L. Tobler, Ductile Fracture with Serrations in AISI 310S 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 .

,

American Society for Testing and Materials, Philadelphia (1983).

6. M. Shimada, R. Ogawa, T. Moriyama, and T. Horiuchi, Development of a Cryogenic Fracture Toughness Test System, in: Cryogenic Engineering,

21:269-274 (1986).

7. H. Nakajima, K. Yoshida, K. Okuno , M. Oshikiri, E. Tada,

S. Shimamoto, R. Miura, M. Shimada, S. Tone, K. Suemune , T. Sakamoto,

and K. Nohara, Fracture Toughness of Newly Developed Structural

Materials for Superconducting Coils of Fusion Experimental Reactor,

in: Advances in Cryogenic Engineering 32:347—354 (1986).

328
EFFECTS OF SPECIMEN SIZE, SIDE GROOVING, AND PRECRACKING TEMPERATURE

ON J- INTEGRAL TEST RESULTS FOR AISI 316LN AT 4 K*

T. Ogata, K. Ishikawa, and T. Yuri

Superconducting and Cryogenic Materials Research Group

National Research Institute for Metals

Ibaraki , Japan

R. L. Tobler, P. T. Purtscher, and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

T. Shoji, K. Nakano, and H. Takahashi

Research Institute for Strength and Fracture of Materials

Tohoku University

Sendai, Japan

The effects of specimen size, side grooving, and fatigue

precracking temperature on J- integral fracture toughness test

results were investigated to support the development of a standard 4-K test procedure. The test material was an AISI 316LN

alloy. The tests used compact specimens and a computer

controlled unloading compliance technique. Fatigue precracking at various temperatures (293, 77, or 4 K) had no apparent

effect on fracture- toughness measurements at 4 K. For plain

specimens, the fracture toughness decreased with increasing

specimen thickness in the range 25 to 38 nun; for side -grooved

specimens, the toughness decreased more.

INTRODUCTION

The single -specimen unloading compliance testing technique 1 * 3 is a useful method to evaluate fracture toughness of materials, especially at

liquid-helium temperature. However, there are several problems, such as the effects of specimen size, side grooves, and regression line, that

should be understood before a 4-K test standard can be established. At

this cryogenic temperature, additional problems occur: serration phenomena, specimen heating, and differences in plastic-zone size. We performed a series of experiments to study the effects of specimen size,

side grooves, and fatigue -precracking temperature on fracture - toughness

test results at 4 K.

*In Advances in Cryogenic Engineering — Materials , vol. 34, Plenum,

New York, 1988, pp. 259—266.

329
MATERIAL AND SPECIMEN DESIGN

The test material was a commercial heat of AISI 316LN stainless

steel in the form of a 54-mm-thick plate in the mill-annealed condition

(solution- treated at a minimum temperature of 1313 K) . Table 1 shows the

chemical composition. The tensile properties as a function of temperature are summarized in Table 2.

Table 1. Chemical Composition for AISI 316LN Steel (wt.%)

c Mn P S Si Cr Ni Cu Mo Co N Fe

0.016 1.10 0.020 0.015 0.54 17.97 10.51 0.17 2.12 0.25 0.139 bal

,

Table 2. Tensile Properties for AISI 316LN Steel

Temperature

(K)

Yield

Strength

(MPa)

Ultimate

Strength

(MPa)

Young's

Modulus

(GPa)

Elongation

(%)

Reduction

of Area

(%)

293 275 580 205 60 68

76 760 1428 205 53 65

4 894 1640 205 44 48

Three sizes of compact (CT) specimens were machined from the plate

in the TL orientation according to ASTM Method E 399-83. Their designations and dimensions are given in Figure 1. The extensometer was attached along the load line, as required in ASTM E 813-81. Modified CT

specimens with the planar dimensions of a 1TCT specimen, but only 12.5 mm

thick, were also used in some preliminary experiments to study the effect

of crack length.

PROCEDURE

Fatigue Cracking

A servohydraulic testing machine was used for fatigue cracking and

Jj c tests. Fatigue cracking was done with a minimum—maximum load ratio

of 0.1 and a frequency of 10 Hz. The last 2.6 mm of the fatigue cracks

were introduced in each size specimen at 293, 77, and 4 K. The maximum

stress intensity was kept low, about 24 MPa°m^, to reduce the plasticzone size. For the 0.5TCT and 1TCT specimens, the relative crack- length

ratio (a/W) was 0.6. For the 1.5TCT specimens, a ratio of 0.7 was adopted to enable testing in a 10-metric- ton-capacity machine.

330
a) 0.5TCT

b) 1 TCT

cl 1.5TCT

1 *-“-

—

<1

-1

1 . —

31.25 125 I 25

12

Figure 1. Specimens used in this study (all dimensions are in mm). The

notations 0.5TCT, 1TCT, and 1.5TCT refer to specimens 12.5 mm,

25 mm, and 38 mm thick, respectively.

After precracking, for two specimens of each size, side grooves were

machined to a net thickness reduction of 25%. The notch-root radii of

the 0.5TCT, 1TCT, and 1.5TCT specimens were 0.25, 0.5, and 0.75 mm, re- spectively. When testing side -grooved specimens, compliance was calculated using the effective thickness, Be ff, and Jjc values were calculated

using the net thickness, as indicated in ASTM E 813-81, sections 8.7.1

and 9 . 1 . 1

.

3 Beff

is given by

Beff - B - (B - Bn) 2 /B , (1)

where B is specimen thickness and Bn is net thickness.

Fracture-Toughness Testing

The standard CT specimens were tested in liquid helium according

to the ASTM Method E 813-81 procedure with a computer-controlled servo- hydraulic testing machine. Figure 2 illustrates the flow chart of this

test. Table 3 shows the initial basic testing conditions. The loading

rate during strain control was 0.4 mm/min. The compliance during un- loading was calculated to within 1%; it includes a rotation correction. 4 The modified CT specimens were tested according to a key-curve method

331
Start

1 Initialize

Loading

,oad-control

Sampl ing

( p , d ) No^| Yes

Unloading

routine —<Td- cr / E>0^

No

Yes

Load inc

Stra in-control

(

Unload ing

) routine J

Stop loading

Hold 20 s I Load control

Start unloa ding

Sampling

(p. d)

Stop unloading

Start re-loading

Stop re-loading

Calculate & d i s p 1 a v C , a , J leturn )

Figure 2. Flow chart of the computer-controlled Jjc test. The symbols

p, d, a, C, and E represent load, displacement, crack length,

compliance, and Young's modulus, respectively.

Table 3. Initial Jj c Test Conditions

0 . 5TCT 1TCT 1 . 5TCT

First unloading (kN) 5[5] 4Q[25] 50 [ 30

]

Unloading ratio 0.2 0.2 0.2

Loading step (kN) 100 200 200

Strain step (mm) 0.05 0.05 0.05

NOTE: Values in brackets are for specimens with side grooves.

332
(described in detail in a separate report 5 ); in these tests, a fracturetoughness parameter designated Jj_ was determined and used to evaluate the

effect of the a/W ratio on the fracture toughness.

After testing, each specimen was heat-tinted (ASTM E 813-81, section

8.3.1) at 593 K for 1 h to mark the final crack length prior to pulling

the specimen apart. The fracture surfaces were examined by scanning

electron microscopy (SEM) for any microscopic effect of the fatigueprecracking temperature on the fracture appearance, particularly at the

location of the original crack tip prior to fracture- toughness testing.

RESULTS AND DISCUSSION

Test results are summarized in Table 4, and typical load—displacement curves for plain (i.e., non-side-grooved) CT specimens are shown

in Figure 3. Large "pop-in" fractures, such as those seen in Figure 3b,

occurred sporadically throughout the testing. Several tests were terminated prematurely (i.e., before a four-point resistance curve could be

obtained), owing to such pop-ins, as noted in Table 4. Figure 4 shows

two examples of the resistance curves. The intersection of the blunting

and tearing lines determines the critical value of J, according to ASTM E 813. 3 The serrated yield behavior is commonplace at 4 K but not recognized by the existing test standard, which is written for a test temperature of 295 K. Large pop- ins due to fast fracture are, in fact, excluded

by the existing standard (see note 2 of ASTM E 813-81). Therefore, the

Jj c values in Table 4 are designated Jq rather than Jj c

. The estimates

of Kj c (J) in Table 4 were obtained from the value of Jq, according to

ASTM E 813-81, section 9.4.

The effects of the test variables (specimen size, fatigue-precracking

temperature, and side grooving) on Kj c (J) are illustrated by Figure 5.

(Invalid data due to skewed crack fronts were also plotted to estimate

effects.) The toughness of specimens decreased with increasing specimen

thickness (at least in the 25- to 38-mm interval), and the toughness of

the side-grooved specimens was lower than that of the plain specimens.

The measured values for the two types of CT specimens, however, were

nearly equal for the largest specimens, 1.5CT. The effects of specimen

thickness and side grooving were attributed to the constraint exerted at

the crack tip

.

There are no clear effects of fatigue-precracking temperature on the

measured toughness in the plain specimens. Fractographs (SEIM) from the

1TCT specimens are shown in Figure 6. The crack front of the specimen

fatigued at 4 K is clearer than that at 293 K and similar to that at

77 K. None of these specimens exhibited stretch zones between the fatigue precrack and the Jj c -test fracture surface. The Jj c -test fracture

surface of each specimen was composed of a uniform distribution of small

dimples (1 to 3 /im) with occasional delaminations.

333
Table 4. Fatigue Testing Matrix and Jj c Results for AISI 316LN Steel

Soecimen

Type Number

Side

Grooves

a/W

Fatigue

Temperature

(K)

^max . (MPa -it 5

)

J Ic

(kJ/m2

) KIc

(MPa-mh

)

0.5TCT Nl-1 yes 0.66 77 22 107 148*t

Nl-2 yes 0.68 77 23 125 160*

Nl-4 no NA 77 23 NA NA*

Nl-5 no 0.58 77 23 160 182*

1 . OTCT N2-

1

yes 0.61 77 24 117 NA It N2-2 yes NA 77 24 134 166*

N2-3 no 0.61 77 24 178 191*

N2-4 no 0.62 77 24 180 192*

N2-5 no 0.61 293 24 235 219**

N2-6 no 0.61 293 24 176 1 90 * N2-7 no 0.61 4 24 189 197*t

N2-8 no 0.60 4 24 217 naII *

1 . 5TCT N3-1 yes 0.72 77 23 119 156*

N3-2 yes 0.72 77 23 112 152 * N3-3 no 0.71 77 25 123 159

N3-4 no 0.71 77 23 129 163*

invalid due to skewed crack front

t "pop-in" fast fracture; test terminated (see Fig. 3b)

$ failure due to machine malfunction

| incomplete resistance curve; Jq is an estimate from less than four

J—Aa points

(a) (b)

Figure 3. Load—displacement curves for (a) 1.5TCT and (b) 1TCT specimens; fast fracture terminated the test of the 1TCT specimen.

334
Figure 4. J-Aa curves for 1 . 5TCT specimens; (a) plain, (b) with side

grooves. (The blunting, tearing, and exclusive lines were drawn according to ASTM E 813. 3

)

335
0.5T 1 T 1.5T

Figure 5. Kj c(J) estimates at 4 K as a function of thickness for proportionally sized specimens (W/B - 2) , showing the effects of

side grooves. Symbols indicate fatigue-precracking temperatures. Open symbols indicate plain specimens; closed symbols,

specimens with side grooves.

The fatigue -precracking temperature had a large effect on the size

of the plastic zone (R) in the specimen at the end of the precracking.

The size of R can be estimated from the loading conditions and the yield

strength (ay) by

R - (K/ay) 2 /6?r (2)

where K is the stress -intensity factor at the end of the precracking.

Using K - 24 MPa-m^, we estimated R to be 0.4 mm at 293 K and 0.04 mm at

4 K. Despite the factor of 10 difference in these values, there was no distinguishable effect of precracking temperature on the measured toughness or the fracture appearance.

Prior to the main series of unloading compliance tests, we briefly

considered the possibility of an a/W effect at 4 K. Performing one test

at a/W ratios of 0.4, 0.5, 0.6, 0.7, and 0.8, we obtained - 179 kJ/m2

, ±6%. This scatter is less than that observed for Jj c tests of six 1TCT

specimens by the compliance technique (Table 4). It is clear that

values are not equal to Jj c because of the different test methodology

used. 3 ’ 5 Regarding the a/W effect, however, the values were internally consistent and therefore indicated a/W independence.

Our results for this heat of AISI 316LN indicate that some of the

restrictions in the ASTM E 813-81 might be unnecessary for the fracture

toughness test standard at liquid-helium temperature. The fatigue crack

336
length and precracking conditions apparently do not affect the measured

toughness significantly. Also, even though all the specimens tested at

4 K met the ASTM E 813-81 thickness requirement (section 9. 3. 5.1, B >

25 Jic/ay) , the results were affected by specimen thickness and side

grooving. Other results for austenitic stainless steels of higher tough- ness than our 316LN material showed that the Jj c values for plain and

side-grooved 1TCT specimens were equal. 6 ’ 7 This contrasts with our re- sults for the AISI 316LN material and suggests that the effect of side

grooving may be material dependent, perhaps because the materials differ

in sensitivity to the constraint to plastic deformation imposed by the

side grooves.

Figure 6. Fractographs (SEM) of a plain 1TCT specimen. (a) and (b) -

fatigue at 293 K; (c) and (d) - fatigue at 4 K. Regions

within the rectangles in (a) and (c) are shown magnified

in (b) and (d)

.

337
CONCLUSIONS

We determined the fracture toughness of specimens of different sizes

at 4 K using the unloading compliance technique. The specimens fractured

in the elastic—plastic region. Trends obtained from this study were

1.

The fracture toughness of plain specimens decreased with increasing

specimen size in the range 1TCT to 1 . 5TCT (thickness range 25 to

38 mm)

.

2.

The fracture toughness of this steel at 4 K is reduced by side

grooves; however, the effect is small, less than 10%, in the largest

specimen size (1.5TCT, or 38 mm thick).

3. Fatigue precracking at 293, 77, or 4 K at a stress -intensity factor

of about 24 MPa-m^ has no measurable effect on the fracture toughness

measured at 4 K.

ACKNOWLEDGMENTS

This work was performed as part of the U.S.—Japan Cooperation

Program in Magnetic Fusion Energy. In Japan, this program is sponsored

by the Ministry of Education of Japan. In the United States, it is

sponsored by the U.S. Department of Energy. We thank K. Nagai at the

National Research Institute for Metals for his contribution to the Jj c test programming. We are also grateful to L. T. Summers of Lawrence

Livermore National Laboratory supplied the test material.

REFERENCES

1. A. Saxena and S. J. Hudak, Jr., Review and extension of compliance

information for common crack growth specimens, Int. J. Fract . 14:453-468 (1978).

2. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, S ingle

specimen tests for Jj c determination, in: "Mechanics of Crack

Growth," ASTM STP 590, American Society for Testing and Materials,

Philadelphia (1976), pp. 27-42.

3. Standard test method for Jj c , a measure of fracture toughness, ASTM

Designation E 813-81, in: 1986 Annual Book of ASTM Standards , section 3, Vol. 03.01, American Society for Testing and Materials,

Philadelphia (1986) pp. 768—786.

4. H. Takahashi, M. A. Khan, and M. Suzuki, J. Test, and Eval . 8:63—67

(1986)

.

5. R. L. Tobler, H. Takahashi, and T. Shoji, Key curve method of fracture toughness determination for austenitic stainless steels at

4 K, in preparation.

338
6. F. J. Loss, B. H. Menke , and R. A. Gray, Jr., Development of J—

R

curve procedures, in: "NRL—EPRI Research Program (RP 886.2),

Evaluation and Prediction of Neutron Embrittlement in Reactor

Pressure Vessel Materials," Annual Progress Report for CY 1978,

J.R. Hawthorne, ed. , NRL Report 8327, Naval Research Laboratory,

Washington, D.C. (August 1979).

7. M. Shimada, R. L. Tobler, T. Shoji, and H. Takahashi, Size, sidegrooving, and fatigue precracking effects on J- integral test

results for SUS 304 stainless steel at 4 K, in: "Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures —XI," NBSIR 88-3082, National Bureau of Standards, Boulder,

Colorado (1988) pp . 317—328.

339
.

.
MECHANICAL TESTS OF LARGE SPECIMENS AT 4 K: FACILITIES AND RESULTS*

K. Yoshida, H. Nakajima, M. Oshikiri,

R. L. Tobler.t and S. Shimamoto

Japan Atomic Energy Research Institute

Ibaraki, Japan

R. Miura and J. Ishizaka

Japan Steel Works

Tokyo, Japan

Mechanical tests of large specimens at 4 K were performed

using a hundred-metric- ton cryogenic hydraulic test machine at

the Japan Atomic Energy Research Institute's Naka Fusion Research Establishment. Specimens were machined from a 140-mm

forged steel plate and from a 100-mm hot-rolled steel plate.

In 4-K tensile tests, data for a 25 -mm- diameter round specimen were obtained and compared with data for a standard 7-mmdiameter specimen. In 4-K fracture tests, data for 75- and

100 -mm- thick compact specimens were compared with data for

25-mm- and 12 . 5-mm-thick specimens of similar geometries. The

large cryogenic hydraulic tensile machine, a ten-metric- ton

cryogenic multispecimen tensile machine, and a ten-metric- ton

cryogenic hydraulic machine are described.

INTRODUCTION

The development of new Japanese cryogenic steels (JCS) , a cooperative effort of the Japan Atomic Energy Research Institute (JAERI) and

several Japanese steel companies that began in 1981, is nearly completed.

Four kinds of steel alloy base metals have satisfied JAERI 's requirements: yield strength > 1200 MPa and fracture toughness > 200 MPam* at

4 K. 1 The gas - tungsten arc weldments of these steels also have good

mechanical properties. During the JCS developmental work, the Demonstration Poloidal Coil (DPC) project 3 required and used JCS in superconducting coils and other cryogenic applications. Details of these DPC

Published in Advances in Cryogenic Engineering — Materials , vol. 34,

Plenum, New York, 1988, pp. 225-232.

^Visiting scientist, from the National Bureau of Standards, Boulder

Colorado, U.S.A.

341
applications (mainly of thin plate) were reported previously. 2 Plates

100 mm or more thick are now required for the toroidal coil cases and

supporting structures of the Fusion Experimental Reactor. For heavysection applications, it is necessary to show that the mechanical properties of the materials are not significantly affected by size effects

arising from the scale-up of designs. Therefore, a new hundred-metric

ton cryogenic tensile machine was installed at JAERI—NAKA in 1986 to

enable very large specimens to be tested. With this machine, the effects

of specimen size on the tensile and fracture toughness behavior of the

structural steels were studied.

TESTING MACHINES

Machines for testing at cryogenic temperatures are indispensable

tools for the development of high-performance superconducting coils.

Innovative technological devices, such as superconducting coils, use materials to the limits of their capabilities. Three mechanical testing

machines have been installed at JAERI, as described in Table 1. These

testing machines are intended not only for steel developmental work but

also for superconductor material and nonmetallic material evaluations.

The ten- ton cryogenic tensile machine (TS10) 4 has both a turret-disk

grip (10 samples per cool-down) and a conventional grip. In this study, we used the turret-disk grip, in which can be installed ten round specimens up to 7 mm in diameter or ten plate specimens up to 4 mm thick. A

strain gage was epoxied to each specimen to measure the yield strengths.

The ten- ton cryogenic cyclic machine (PL10) was designed for fracture toughness (Jjc ) , fatigue crack growth rate, and conventional fatigue

data measurements. It was used in this study to determine the fracture

toughness of the relatively small or standard compact specimens, as shown

in Table 1.

Table 1 . Machines at JAERI for Testing Structural Materials at 4 K

Machine TS10 PL10 TS100

Capacity (metric ton) 10 10 100

Feature turret disk fatigue test large size

Specimen

Tensile — Round (mm in diam.) 7 7 25 - Plate (mm thick) <4 NA NA

Compact tension coding IT 0.5T, IT 3T, 4T

Number of specimens

tested per day 10 1 0.3

Consumption of liquid

helium per specimen (i) 8 80 800

342
Fig. 1. The TS100 (A) and PL10 (B) cryogenic hydraulic test machines.

The large cryogenic tensile machine (TS100) shown in Fig. 1 was designed to determine the tensile properties and the fracture toughness

of steel or superconductor specimens under monotonic loads of up to 100

tons. The total height of the TS100 is about 5 m, and the height of the

section to be immersed in cryogenic liquids is 3 m. Figure 2 shows the

installation features and grips for a fracture toughness specimen. The

shear pins were machined from the same JCS base material because high

yield strength is required to resist bending. The space for testing in

the TS100 is 1.8 m long and 0.8 m in diameter. This large testing space

will be useful in performing several verification tests in the future

343
The cryogenic apparatus and the hydraulic cylinder are assembled as one piece; before testing, this assembly is lifted by crane to its position

on the test stand. The cryostat is then jacked up and its flange is

fastened with bolts. The time required for one measurement cycle (cool

down, testing, and warm-up) is three days. The helium consumed during

one cycle of operation is about 800 Z, but the heat load to the liquid

helium is less than 2 W.

780

Fig. 2. The TS100 cryogenic apparatus with

a compact specimen installed.

(All dimensions are in mm)

.

344
SPECIMENS

Two JCS materials were selected for this study:

1. CSUS—JJ1: 5 lCr—12Ni—lOMn—5Mo—0 . 2N, 140-mm forged plate

(from a 6-metric-ton heat, electroslag remelted)

2. CSUS—JNl

6 25Cr—15Ni—4Mn—0 . 35N, 100-mm hot-rolled plate

(from a 50-metric- ton heat)

For tensile tests, 7- and 25 -mm- diameter round specimens (as

shown in Fig. 3) were machined from the CSUS-JJ1 plate in accordance

with Japanese Industrial Standard JIS Z 2201. 7 The specimens were

taken from the plate at midthickness and were axially oriented in the

long transverse direction. The reduced section lengths of the 7- and

25-mm-diameter specimens were 42.4 and 175 mm, respectively. The

specimen ends were threaded for gripping, as shown in Fig. 3.

For fracture toughness tests, CT specimens were machined from

the two materials, in accordance with ASTM Method E 813-81. 8 The CT

specimens, as shown in Fig. 4, were proportionally sized with thicknesses of 12.5, 25, 75, or 100 mm; the coding used in this study is

indicated in the figure caption.

Fig. 3. Tensile specimens machined from the CSUS plate.

The large one is 25 mm in diameter; the small

one is 7 mm in diameter.

345
Fig. 4. Compact specimens: 100 -mm thick (4-T CT) , 75-mm thick (3-T CT)

,

25 -mm thick (1-T CT) , and 12 . 5 -mm thick (0.5-T CT)

.

TEST RESULTS

Tensile Tests

The results of tensile tests at 4 K are shown in Table 2 and Fig. 5.

The testing speeds were 0.6 mm/min up to the yield strength and 3 mm/min

up to fracture. The nominal strain rate for the 25 -mm specimen was about

one -fifth that of the 7 -mm specimen because the large specimen had a longer reduced section length. The effects of size on the tensile properties were

1. The yield strengths (ay) and elongations (El.) were nearly the same for both specimen sizes.

2. The ultimate strength (au) and reduction of area (R.A.) were 4 to 8%

less for the 25-mm-diameter specimen than for the 7-mm-diameter

specimen.

3. The post- test magnetizations for both specimen sizes were negligibly

small

.

From a design viewpoint, the differences in the ultimate strength

and reduction of area are negligibly small: the size of the 7- and 25-mmdiameter specimens did not affect the tensile properties significantly.

Our results indicate that data from small (7-mm) specimens are sufficient

to characterize these cryogenic structural steels.

346
Table 2. Results of JJ1 Tensile Tests at 4 K

Specimen

Diameter

(mm)

0.2% ay

(MPa) (MPa)

El.

(%)

R. A.

(%)

Relative

Magnetic

Permeability

25 1142 1491 43.2 46.1 1.018

7 1132 1574 44.7 54.3 1.006

1127 1570 43.6 50.0 1.009

I

Fig. 5. Specimen size effects on tensile properties at 4 K.

Fracture Toughness Tests

The fracture toughness test results are summarized in Table 3 and

Fig. 6. The Kj c values were calculated from the Jj c values according to a formula given in section 9.4 of ASTM E 813-81. 8 The effects of specimen size on fracture toughness were

1. For the CSUS-JN1 steel, there was no size effect on fracture tough- ness for the range of specimen sizes studied.

2. For the CSUS-JJ1 steel, the 0.5-T CT specimen had a Kj c (J) fracture

toughness that was 10% larger than those of the 1-T CT and 4-T CT

specimens

.

3. No unstable linear-elastic failures were observed in any of the

tests, including the tests of the large 3-T CT and 4-T CT specimens

347
There was no significant difference in fracture toughness measurements for different specimens of the CSUS-JN1 steel, but the fracture

toughness values for the 0.5-T CT specimens of CSUS-JJ1 were larger than

those for the larger specimens of that steel. From a design viewpoint,

the difference in values for the 1-T and 4-T CT CSUS-JJ1 specimens is

negligibly small. Therefore, from the data presently available, we con- clude that conventional laboratory tests performed using the usual 1-T CT

specimen size are sufficient to characterize the fracture toughness of

these steels at 4 K. We intend to examine the fracture surfaces microscopically and to confirm the trend of these results by future tests of

other materials.

Table 3. Results of Fracture Toughness Tests at 4 K

Material JN1 JJ1

Specimen J Ic

(kJ/m2

) KlcW

(MPa -in 4

) Specimen J Ic

(kJ/m2

) KIC (J)

(MPa ‘ni

4

)

3 = T CT 205 203 4-T CT 343 257

218 210

1-T CT 222 216 1-T CT 375 269

213 208 369 266

0.5-T CT 203 203 0.5-T CT 425 286

225 213 475 302

449 294

Fig. 6. Specimen size effects on fracture toughness, Kj c (J), at 4 K.

348
DISCUSSION

As shown in Table 3, the yield strengths of the thick plates tested

in this study are lower than the 1200 MPa value specified by JAZRI for

thinner plates during the JCS development. Therefore, further considera

tion will be given to this requirement with respect to thick plates.

The fracture toughness requirement specified by JAZRI 1 was determined from the design concept that a leak will occur before failure in

structural applications. When, under 500-MPa- amplitude stress cycles, a initial crack grows through the thickness of the plate, the stress inter,

sity factor (K) is 200 MPa -in’1. The CT- specimen tests simulated this con

dition, and yet the thick 3-T CT and 4-T CT specimens of CSUS-JN1 and

CSUS-JJ1 did not completely fracture. Evidently the CSUS-JN1 and CSUS

JJ1 steels have larger mechanical safety factors than we anticipated;

they are now considered ready to be used in coil cases and support struc

tures (more than 100 mm thick) of superconducting coils for fusion reactors. Other JCS steels will be tested in the TS100 in the near future.

CONCLUSIONS

Large specimens were tested using JAZRI 's hundred-metric- ton

cryogenic hydraulic tensile machine. The principal results of the effec

of size on tensile properties and fracture toughness were

1. Specimen size had little or no effect on the tensile test results of

7- and 25 -mm- diameter round specimens.

2. The fracture toughness varied little from the 1-T CT to the 4-T CT

specimens

.

3. The fracture toughness values of the 0.5-T CT specimens of CSUS-JJ1

were slightly higher (about 10%) than that of the 4-T CT specimen.

4. No linear-elastic fractures were observed in the 3-T CT and 4-T CT

specimens of CSUS—JN1 and CSUS—JJ1 steels, respectively.

ACKNOWLEDGMENTS

The authors thank Drs . S. Mori, K. Tomabechi, and M. Tanaka for

their continuous encouragement during this study. The contribution of

the Nippon Steel Company is also hereby acknowledged. The study was partially supported by the U.S. Department of Energy, Office of Fusion

Energy.

349
REFERENCES

1. K. Yoshida, H. Nakajima, K. Koizumi, M. Shimada, Y. Sanada,

Y. Takahashi, E. Tada, H. Tsuji, and S. Shimamoto, Development of

cryogenic structural materials for tokamak reactor, in: "Austenitic

Steels at Low Temperatures," R. P. Reed and T. Horiuchi, eds . , Plenum

Press, New York (1983), pp. 29—39.

2. H. Nakajima, K. Yoshida, M. Oshikiri, T. Takahashi, K. Koizumi,

S. Shimamoto, M. Shimada, S. Tone, S. Sakamoto, K. Suemune, and

K. Nohara, Tensile properties of new cryogenic steels as conduit

materials of forced flow superconductors at 4 K, in: "Advances in

Cryogenic Engineering—Materials," vol. 34, Plenum Press, New York

(1988), pp. 173-181.

3. H. Tsuji, Design selection for the fabrication on the demonstration

poloidal coil, in: "Proceedings, 7th ANS Topical Meeting on the

Technology of Fusion Energy," June 15—19, Reno, Nevada (1986).

4. K. Yoshida, Y. Takahashi, E. Tada, M. Shimada, and S. Shimamoto,

Mechanical properties of the Japanese LCT coil conductor, in:

"Advances in Cryogenic Engineering-Materials," vol. 28, Plenum

Press, New York (1982), pp. 781-790.

5. H. Nakajima, K. Yoshida, K. Okuno, M. Oshikiri, E. Tada,

S. Shimamoto, R. Miura, M. Shimada, S. Tone, K. Suemune, T. Sakamoto,

and K. Nohara, Fracture toughness of newly developed structural

materials for superconducting coils of the fusion experimental

reactor, in: "Advances in Cryogenic Engineering—Materials ,

vol. 32, Plenum Press, New York (1986), pp. 347—354.

6. K. Suemune, T. Sakamoto, T. Ogawa, T. Okazaki, and S. Maehara, The

manufacture and properties of nitrogen- containing Cr-Mn and Cr-Ni

austenitic stainless steels for cryogenic use, in: "Advances in

Cryogenic Engineering—Materials," vol. 34, Plenum Press, New York

(1988), pp. 123-129.

7. Japanese Industrial Standard JIS Z 2201 in: "JXS Handbook, Steel1981," p. 23.

8. Standard Test Method for Jjc , A Measure of Fracture Toughness, ASTM

E 813-81, in: "1986 Annual Book of ASTM Standards," American Society

for Testing and Materials, Philadelphia (1986), pp. 768-786.

350
ACOUSTIC EMISSION AND ITS APPLICATIONS TO FRACTURE STUDIES

OF AUSTENITIC STAINLESS STEELS AT 4 K*

H . Takahashi and T . Shoj

i

Research Institute for Strength and Fracture of Materials

Tohoku University, Sendai, Japan

R. L. Tobler

National Bureau of Standards

Boulder, Colorado

In contrast to their silent cracking behavior at room

temperature, the austenitic stainless steels are highly

active acoustically during deformation and fracture at 4 K.

Acoustic emission is promoted by low- temperature strengthening and by discontinuous plastic flow in ductile alloys near

absolute zero. Consequently, the prospects are quite good

for acoustic-emission monitoring as a tool for the study of

austenitic steels at 4 K. This paper illustrates some potential applications.

INTRODUCTION

Acoustic emission (AE) may be defined as energy in the form of elastic waves generated by solids undergoing abrupt changes in stress. For

three decades, acoustic emission at room temperature has been employed to

study materials, to evaluate structures, and to monitor processing operations. 1 * 2 Cryogenic applications are also growing. Acoustic emission

has been used to monitor the operation of superconducting coils and magnet assemblies 3 * 4 and to study the microdeformation and fracture of

epoxies, 4 * 5 composites, 6 and pure copper. 7

At Tohoku University's Research Institute for Strength and Fracture of Materials (RISFM) , AE techniques are often used to characterize

structural steels at room temperature. During conventional fracture

mechanics tests, acoustic emission monitoring can provide detailed information about alloy microfailure mechanisms and fracture processes

while the usual macroscopic fracture parameters are being determined.

In Advances in Cryogenic Engineering — Materials

,

vol. 34, Plenum,

New York, 1988, pp. 387-395.

351
Recently, the same techniques were applied to austenitic stainless

steels at 4 K. 8-10 Selected results for three alloys are reviewed here

to illustrate possible applications.

MATERIALS

The materials tested were two wrought and annealed base metals

(Fe—18Cr—3Ni—13Mn and AISI 310S BM) and a shielded metal arc weld (AISI

310SWM). 9

’ 10 Table 1 summarizes the relevant material properties and

fracture characteristics. Comparison in Fig. 1 of the strength—toughness

data with established trends shows that the three studied materials are indeed representative commercial alloys: Fe—18Cr—3Ni—13Mn and AISI 310S

BM fall near the low- and high- toughness extremes for base metals; 11

AISI 310S WM, with moderate fracture toughness, lies at the center of

the trend for 316- and 308-type welds. 12 Representative scanningelectron microscope (SEM) fractographs from compact specimens tested

at 4 K are shown in Fig. 2.

Table 1. Fracture characteristics of test materials at 4 K

Material

Curve

Type

ay

/CTu

( ~ )

J Ic

(kJ/m2

)

Kic

(MPa-m^)

Fracture

Mechanism

Dimple

Size

(/jm)

Fe—18Cr—3Ni—13Mn LE 0.85 24a 71 slip -band cracking 0

AISI 310S WM EP 0.78 118 159b void coalescence 4

AISI 310S BM FP 0.60 333 293b void coalescence 25

Notes: a estimated from Kj c measurement; b estimated from Jj c measurements.

METHODS

Tests were conducted with compact specimens (22 or 25 mm thick, TL

orientation9 ’ 10

) submerged in liquid helium at 4 K. A conventional system for detecting, conditioning, and analyzing AE signals is shown in

Fig. 3. Such a system 13 was modified for tests at 4 K by employing a waveguide and reducing the gain. The waveguide was a 4-mm-diameter

stainless steel rod, welded to the back face of the specimen. The guide

transmitted AE signals from the specimen at 4 K to a wideband piezoelectric transducer at 295 K via a coupling outside the cryostat. A significant problem at 4 K is signal saturation owing to the strong emissions

from these steels. Therefore, a 20-dB gain was preferred to the usual

60-dB gain at 295 K. The signals were fed into a distribution analyzer,

which counted pulses according to their peak amplitudes and determined

the individual event energies, E^jr. The event energy is the square of

352
Fig. 1. Strength—toughness data for steels used

in the acoustic emission studies.

the output, V, in volts, integrated over the period of its duration.

Sometimes the peak volt-age , Vp, or its square is taken as a relative

measure of signal energy. The square of the peak voltage is proportional to, and correlates well with, the conventional measure of energy.

ACOUSTIC ACTIVITY

Not all materials emit useful AE signals. According to the acoustic rating of steels at room temperature by Ingham et al., annealed

austenitic stainless steels rank last. 14 The low yield strengths, high

ductilities, high strain-hardening capabilities, and stable cracking

mechanisms that typify austenitic steels at 295 K are just not conducive

to abrupt physical processes that produce acoustic emission. At - K.

the circumstances are changed: all the austenitic steels are significantly stronger, their ratios of yield strength (av ) to ultimate tensile

strength (<7U) are larger, and the discontinuous deformation process

becomes prominent in the ductile alloys. Under these circumstances

significant emissions are likely.

Our survey of austenitic stainless steels includes one representative for each of the three categories: linear elastic (LI) , elasticplastic (EP) , and fully plastic (F?) . Load—displacement ;?—6 ) records

and acoustic-emission outputs for each case are shown in Fig.

353
Fig. 2. Fractographs : A) large dimples, AISI 310S BM;

B) small dimples, AISI 310S WM;

C) facets, Fe—18Cr—3Ni—13Mn steel.

Markers are 70, 20, and 2.5 pm, respectively.

The P—5 curve for the high-strength Fe—18Cr—3Ni—13Mn steel is nearly

linear except for small deviations from minor pop -ins near the maximum

load. The AE begins at low loads and continues until fast fracture

finishes the test with a bang.

The AISI 310S weld- and base-metal steels are both ductile at 4 K,

and both display nonlinear P—5 curves with serrations. The base metal

has much higher toughness and smaller, more numerous serrations; its

fully plastic behavior is distinguished from the weld's elastic—plastic

behavior by its approach to a limit load. For these materials, acoustic

354
Fig. 3. Acoustic- emission monitoring apparatus.

activity increases as nonlinear loading develops. The event energies

increase in magnitude at higher displacements. Later, serrations begin

and load drops occur with increasing magnitude accompanied by high-level

AE bursts.

In summary, the austenitic stainless steels at 4 K are acoustically

active. Pop -ins and serrations are associated with audible AE, as is

well-known, but acoustic emissions also occur in early stages of loading.

Thus, significant AE precedes macroscopic fracture as measured by the J- c or Kj c parameters. The precursory emissions are associated with various

deformation and cracking processes and may be useful for purposes such as •those described below.

APPLICATIONS

Histograms of Acoustic-Emission Events

A representative histogram of AE events for AISI 310S base metal at

4 K is shown in Fig. 5. Usually, histograms for steels at 295 K show a trend of continuously decreasing event numbers (N) at increasing voltages

according to the relationship N oc Vp^, where b < 0. This trend is basic

evidence of a single-emission mechanism. Instead, the AISI 310S base

metal at 4 K shows a bimodal distribution: two active ranges (V D - 0.20

—0.65 and 1.0 — 4.0 V) separated by a quiet range (Vp - 0.65 — 1.0 V)

This indicates two major sources of acoustic emissions. Thus, auster.:*::

stainless steels are not only highly active at 4 K, but their event

histograms differ markedly from typical room- temperature results

355
m

CM>

UJ

>O

ccUJz

UJH

Z

UJ>

UJ

UJ<

Fig. 4. Load—displacement records and AE outputs for steels at 4 K:

A) linear-elastic, B) elastic—plastic , C) fully plastic behavior.

356

AE

EVENT

ENERGY,

E^e

,
Fig. 5. Events histogram for AISI 310S

base metal at 4 K.

Fractograthv and Identification of Acoustic-Emission Sources

Since fracture in commercial steels occurs as a series of events,

many acoustic sources are possible: inclusion-matrix separation, void

growth and linking, fast shear, and various types of brittle cracking,

to name a few. Following previous studies

,

1

5

* 1

8

one can attempt to

identify the principal AE sources by correlating event energies with

fractographic data from microexaminations. For example, SIM examinations

indicated that secondary cracking in the fracture zone near the fatigue

crack tip was a possible AE source in the AISI 310S base metal. In

Table 2, the number and the lengths of the secondary cracks are compared

with the corresponding AE histogram of Fig. 5. As indicated in the

table, the number of AE events in the range V

p

- 1.0 — 4.0 V nearly

equals the number of long secondary cracks observed. Therefore, sec- ondary cracking might constitute an AE source for the higher emission

levels, but this conclusion is tentative since only one specimen was

examined. The events in the range V

p - 0.30 - 0.65 V do not correlate

with secondary cracking but may be due to other processes, such as abrupt plastic deformation.

Fractography revealed no secondary cracks in the AISI 310S weld

metal. Instead, high-level emissions were associated with pop -in crack

extensions at local areas along the crack front. Pop- ins are increments

of subcritical crack extension, and the term normally implies brittle

cracking. The distinction between brittle and ductile behavior is some- what blurred in the case of this intermediate- toughness veld. Cracking

occurs in rapid spurts, but post-test inspection shows no faceted fracture, only fine dimples 4 /im in diameter. Thus, both the AISI 31 J S case metal and the weld show dimpled fracture surfaces; they also have nearsequivalent yield strengths, but the weld's toughness is compromised,

probably by its high inclusion content.

357
Table 2. Evaluation of secondary cracking as an AE source in AISI 310S base metal at 4 K

Secondary Number of AE Event Number

Crack Length Secondary Level of AE

(mm) Cracks (volts) Events

0 to 0.05 4 0.30 to 0.65 33

0.05 to 0.10 16 0.65 to 1.0 7

0.10 to 0.15 8

1 0.15 to 0.20 0 I

- 10 1.0 to 2.0 11

0.20 to 0.25 2. 0.25 to 0.30 4] X _ c 2.0 to 4.0 7

0.30 to 0.35 2J L 0 /

The Fe—18Cr—3Ni—13Mn material fails by slip-band cracking, a form

of cleavage on {111} planes, as verified by x-ray diffraction and slipline trace analysis. This brittleness is produced by the high nitrogen

content (0.37%). A probable emission source is localized cracking of

individual grains or groups of grains. If so, the microfailure process

for this high strength steel may well parallel that discussed for other

steels of similar strengths that fail by cleavage at room temperature. 18

The microfailure mechanisms surmised for each steel could be verified or refuted by amassing more evidence from a series of specimens

where the tests are interrupted for analysis at progressive stages of

fracture. The procedure is fairly laborious, however, and beyond our present scope, which is to illustrate the principles of emission source identification.

Microfracture Toughness

The cumulative sum of AE event energies is plotted versus the

applied stress - intensity factor, K, on logarithmic coordinates in Fig. 6.

Plots of this type for steels at 295 K demonstrated bilinear trends, with

the slope transition points designated K^£. 15-18 The parameter Kj^jr

marks the point of a rapid increase of emissions and is defined as the

microfracture toughness of the alloy.

The steels at 4 K do not always show simple bilinear trends, but

values can still be identified. As shown in Fig. 6, the data for

two AISI 310S base metals indicate fair agreement. The results for welds

(not shown) were more scattered, reflecting the greater material variability of welds.

Microfracture toughness values and macrofracture toughness measure- ments are compared in Fig. 7. The parameter is a more sensitive,

conservative measure of toughness than Kj c

. The ratio Kiae/Kic decreases

as the level of plastic constraint increases from the FP to the LE case.

358
Fig. 6. Cumulative sums of AE energy output versus applied

stress -intensity factor for AISI 310S base metals.

Fig. 7. Microfracture—macrofracture parameter comparisons

for alloys at three levels of toughness.

359
Only one test for the LE case was completed,, and a very low Kj^£ value

was obtained. The observed low value has a precedent in high-strength

steels at 295 K where the microtoughness values as low as 50% of Kj c have

been reported. 18

Specimen Heating

Adiabatic heating is prominent in cryogenic studies. The specimen

temperature rise, AT, during deformation at 4 K can be measured using

thermocouples; it may also be calculated with assumptions about plastic

work. We previously explored a third possibility, seeking correlations

between AE outputs and temperature rises in compact specimens where AT at

the advancing crack tip was measured by means of thermocouples embedded

in the ligament along the fracture plane. 10

During continuous loading in AISI 310S base- and weld-metal specimens, before the occurrence of serrations, we found a relationship between the rate of acoustic emission and specimen temperature that was dependent on crosshead speed. The change of total AE energy per unit

time versus T/At at crosshead speeds of 0.1 and 1.0 mm/min was monitored.

A threshold was found in the AE rate at about 8,5 x 10" 2 V2

• s/s , below

which AT during continuous loading was zero for these steels. In the

discontinuous yielding regime, the AE energy correlated with temperature

rises for individual serrations. The existence of correlations between

AE and adiabatic heating suggests that these two quantities describe the same physical processes. Thus, future AE measurements could lead to a better understanding of discontinuous plastic deformation at 4 K.

SUMMARY

1. A conventional apparatus employing a waveguide and reduced gain can

be used to monitor acoustic emission during fracture mechanics tests

at 4 K.

2. Austenitic stainless steels of all three toughness levels are acous- tically active during deformation and fracture at 4 K. This fact is

attributed to increased tensile flow strengths and to discontinuous

plastic flow at 4 K.

3. The acoustic emissions precede macroscopic fracture and may be used

to study microfailure mechanisms or monitor the overall fracture

process. Our initial results for three commercial steels indicate

that it is possible to: (1) relate the acoustic emission to specific

micromechanisms for brittle and ductile fracture, (2) define microfracture toughness values at 4 K, and (3) establish correlations

between acoustic emission and adiabatic heating.

360
ACKNOWLEDGMENTS

This work was sponsored in Japan by the Ministry of Education of

Japan and in the United States by the Department of Energy, Office of

Magnetic Fusion Energy. The authors thank Dr. T. Hashida for helpful

discussions

.

REFERENCES

1 .

2 .

3.

4 .

7.

8 .

9.

H. N. G. Wadley, C. B. Scruby, and J. H. Speake. Int. Met. Rev

25:41-64 (1980).

H. N. G. Wadley and R. Mehrabian. Mater. Sci. Fng 6S-9

(1984).

R. W. Fast, ed. , Adv, Crvo. Eng. 31:259-331 (1986).

H. Fujita and Y. Iwasa, Exp. Mech. 26:128-135 (1986).

5. Y. Yasaka and Y. Iwasa, Cryogenics 24:423-428 (1984).

6. T. Nishiura, K. Katagiri, S. Owaki, and T. Okada. Cryogenics

24:329-333 (1984).

H. Fujita, E. Bobrov, and Y. Iwasa, in: Progress in Acoustic

Emission—III , Japanese Society for Nondestructive Inspection Tokyo

222-229 (1986). ’ y ’ K. Ohnishi , Master's Thesis, Research Institute for Strength and

Fracture of Materials, Tohoku University, Sendai, Japan (1986).

K. Ohnishi, T. Shoji, and H. Takahashi, Trans. ISME (A) 52:170—173

( 1986)

.

10. R. L. Tobler, T. Shoji, H. Takahashi, and K. Ohnishi, in: Progress

in Acoustic Emission—III , Japanese Society for Nondestructive

Inspection, Tokyo, 453-461, (1986).

R. L. Tobler, D. T. Read, and R. P. Reed, in: Fracture MechanicsThirteenth Conference , ASTM STP 743, American Society for Testing and Materials, Philadelphia, 250-268 (1981).

R. L. Tobler, T. A. Siewert, and H. I. McHenry. Cryogenics

26:392-395 (1986). —

H. Niitsuma, M. Kikuchi, H. Takahashi, M. Suzuki, and R. Sato, Proceedings , Fifth International Acoustical Emission Symposium

Japanese Society for Nondestructive Inspection, Tokyo, 411—420

(1980)

.

U' 247^293 ’(1975i

SC0 ": ’ C°Uan ' Int ’ J ' Press - Ves ' pi Pi"e

15

' 21-89^99

a (198n'

A ' Khat'' M ' KlkUCh1 ' and M ’ Su2uki ’ Ex°- 16. M. A. Khan, T. Shoji, and H. Takahashi, Res. Mech. Lett 1:133—138

( 1981 ) . 17. M. A. Khan, T. Shoji, and H. Takahashi, J. Test. Eval. 10:3-11

(1982)

.

18. M. A. Khan, T. Shoji, and H. Takahashi, Met . Sci 16:118-126 (1982

11 .

12 .

13.

361
'
U.S. —JAPAN COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

FIFTH DRAFT

SEPTEMBER 1987

NOTE: This is a working document of a standard being developed by the above

committee. Some technical aspects are still being resolved and are subject to revision. The document in its present state is intended

for committee review purposes, but not for general distribution at

this time

.

Proposed Standard Methods for

FRACTURE TOUGHNESS TESTING OF AUSTENITIC STAINLESS STEELS

AT LIQUID HELIUM TEMPERATURE

1 . Scope

1.1 This standard describes the equipment, procedures, and data analysis used for measuring the fracture toughness of structural alloys in liquid

helium [4 or 4.2 K (—452°F)]. Procedures described in ASTM Methods E 399-83

and E 813-81 are adopted and modified to address the specific needs and con- cerns of cryogenic testing.

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, and it is 4.0 K at NBS in Boulder,

Colorado, 1677 m (5500 ft) above sea level. In this standard the test temperature is nominally designated 4 K.

1.2 To measure toughness according to this standard, a fatigue - cracked

specimen is cooled in a cryostat to 4 K and then loaded at a moderate speed

while fully submerged in normal liquid helium (He I) . The standard specimen

is a pin- loaded compact specimen, nominally 25 mm thick and 50 mm wide.

1.3 "This standard specifies the measurement of Kj c , J c , or Jj c , whichever is appropriate for the material behavior at 4 K. For fully plastic

fracture, Jj c is determined by using a resistance curve technique. Measurements of the resistance -curve slope, AJ/Aa, or the material’s tearing

modulus are not included in this standard.

1.4 Values stated in SI units are treated as primary; values stated in

U.S. customary units, secondary.

1.5 This standard does not address safety issues. Safety guide Lines

for handling liquid helium and other cryogens are available elsewhere ( 1

;

It is the user's responsibility to implement appropriate safety practice.-;

363
2. Applicable 2.1 Documents

ASTM Standards :

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 8 Methods of 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 of Tension Testing of Metallic Materials, Metric

( Annual Book of ASTM Standards , Vol. 03.01)

E 399 Standard Test Method for Plane-Strain Fracture Toughness

of Metallic Materials ( Annual Book of ASTM Standards , Vol. 03.01)

E 616 Terminology Relating to Fracture Testing

( Annual Book of ASTM Standards , Vol. 03.01)

E 813 Standard Test Method for Jic> A Measure of Fracture

Toughness ( Annual Book of ASTM Standards

,

Vol. 03.01)

3. Definitions

3.1 The terms relating to tensile testing as defined in Definitions

E 616 shall apply in this standard.

3.2 stress intensity factor, plane- strain fracture toughness , and crack

plane orientation—the definitions in Methods E 399 apply in this standard.

3.3 J- integral , physical crack size, original crack size, crack extension, effective yield strength , blunting line, initiation of stable crack

growth, regression line slope, and normalized crack size—the definitions in

Methods E 813 shall apply in this standard.

3.4 fracture cryostat—an apparatus consisting of a loading frame and

Dewar for testing specimens in cryogenic environments. A schematic is shown

in Fig . 1

.

3.5 Dewar—a vacuum- insulated container for cryogenic fluids.

3.6 adiabatic heating—the internal heating of a specimen caused by

fracture testing under conditions such that the heat generated by plastic

work cannot be quickly dissipated to the cryogen.

3.7 Jc—a critical value of J at the onset of fast fracture causing

large, uncontrolled crack propagation.

364
1 - Dewar pivot f raae 19 - Isolator compression r 2 - Specimen clevis 20 - Adapter bar

3

- Shear pin 21 - Test machine frame

4 - Pivot plate 22 - Load cell

5 - Outer Dewar top 23 - dip gage

6 - Inner Dewar danp 24 - Stainless 3teel Dewar

7 - Transfer hose adapter 25 - Stainless steel Dewar

8 - Adapter plate 26 - J-integral contact ten

9 - Co

i

use adapter fracture specimen

10 - Compression colons: 27 - Seal ring

11 - Clevis pull rod 28 - Liquid heliua level me 12 - Pull rod keeper 29 - Socket head cap screw

13 - Pull rod split ring isolator 30 - Socket head cap screw

14 - Load plate 31 - Socket head cap screw

15 - Pull rod coupling 32 - Socket head cap screw

16 - Load cell adapter 33 - !»ut 17 - Load coluao 34 - Threaded rod

18 - Load cell adapter 35 - Shear pin

Figure 1. Cryogenic fracture- toughness test facilitv

365
4. Summary of Methods

4.1 The object of this standard is to measure the fracture toughness of

austenitic stainless steels at 4 K using a K- or J-based fracture criterion.

The procedures described in Methods E 399-83 and E 813-81 are adopted and

modified in this standard for applications at 4 K. Familiarity with those

Methods is assumed.

4.2 The compact specimen is emphasized here because it offers a larger

fracture- toughness measurement capacity for its size than any other configuration. A relatively low load applied to this specimen produces a large

stress - intensity factor; this enables the mass of the cryostat, grips, and

load train to be minimized, lowering the cryogen consumption and cost per

test

.

4.3 To measure the fracture toughness according to this standard, a fatigue-cracked specimen is cooled to 4 K in a suitable cryostat containing

liquid helium. Using crosshead, stroke, or clip-gage displacement control,

the specimen is slowly loaded while the load-vs-displacement curve is autographically recorded. The type of load-vs-displacement curve observed for

the material in the thickness tested will then determine the appropriate

method of analysis and fracture criterion to be applied (see Fig. 2).

4.3.1 The parameter Kj c is used to characterize materials with a Type I record where the loading is linear-elastic and the crack-tip plastic zone is

small compared with the crack size and specimen dimensions. Except for the

provision of the cryogenic test apparatus, the method of Kj c measurement in

this standard is similar to that set forth in Methods E 399-83.

4.3.2 The parameter J c is used in this standard to characterize Type II

records where the loading is nonlinear and the plastic zone at the crack- tip

is sizable. Although it is not a material constant, J c can be used to characterize the toughness of alloys that lack sufficient section thickness or toughness for valid Kj c or Jj c tests. For tests terminating with a single

fast fracture or a series of large pop-ins (increments of fast fracture), a

J c value may be calculated from the area under the load-vs-displacement

curve up to the critical point of instability. The J c value will vary with

specimen type and size; also, there will be sizable scatter, even for a constant geometry, owing to statistical variation of various factors

influencing unstable fracture.

4.3.3 The parameter Jj c is used to characterize materials displaying a Type III record where the loading is nonlinear (elastic-plastic, or fully

plastic) , where the crack-tip plastic zone is large compared with the crack

size and specimen dimensions, and where a satisfactory resistance curve can be obtained for the material. The recommended method for Jj c measurement at

4 K is the single- specimen unloading-compliance procedure. That procedure

was first developed (2—4) for room temperature tests and was later applied

to austenitic stainless steels at 4 K (5—7)

.

366
I. Linear-Elastic Behavior: K,c

is measured

II. Elastic-Plastic Behavior: J is measured

1. Fully Plastic Behavior: J (c is measured

(sizable plastic zone)

Figure 2. Test records and fracture parameters used at 4 K.

367
4.3.4 A "key- curve" method for characterizing fracture in terms of a parameter has been applied to austenitic steels at 4 K, with some success

(8,9). This method is an estimation procedure rather than a measurement

procedure, since it involves an assumption about the relationship between

crack initiation and crack growth resistances. Since the underlying assumption does not universally hold true for all steels, the parameter Jj_ determined by the key-curve approach is not recognized by this standard at this

time

.

5. Significance and Use

5.1 The significance and use of Kj c as a material toughness parameter

at 4 K is the same as that discussed in Methods E 399-83, section 4.

5.2 The parameter J c defined in this standard is not a true material

constant since it depends on test specimen geometry. The significance of J c values with respect to structural design is not addressed in this standard.

The J c measurement is used here only as a relative measure of toughness for

materials not amenable to Kj c or Jj c characterizations in practical section

thicknesses

.

5.3 The significance of Jj c as a toughness parameter at room temperature

is discussed in Methods E 813-81, section 4. As originally formulated, the

J- integral approach applies strictly to metals that fracture by slow- stable

tearing. At 4 K, ductile fractures typically involve discontinuous plastic

deformation, and the corresponding test records are marked by serrations

(abrupt load drops and momentary temperature rises), as shown in Fig. 3 (10).

Instabilities may also arise from pop- in behavior in brittle materials.

Consequently, resistance curve slope and tearing modulus calculations at 4 K

lack the significance they have at ambient temperatures where crack growth

occurs by a slow-stable, isothermal tearing process. To date, however, the

experimental evidence (11—13) indicates that the occurrence of serrations

does not undermine the use of Jj c as a size- independent material property

for austenitic stainless steels at 4 K.

5.4 The Jj c values obtained for high- toughness alloys may be used to

estimate Kj c values. Conversions from Jj c to Kj c are performed following

Methods E 813-81, section 9.4:

KIc (J) - (Jlc x E) 4

Such conversions are useful for comparing different materials quantitatively

on the basis of a single parameter. The estimated values are identified by

the symbol Kj c (J), to distinguish them from Kj c measurements. Thus, Kj; c (J)

values are obtained using small test specimens that do not display linearelastic behavior in practical section sizes, but the conversion from Jj c to

Kj c (J) is only valid for Type III records.

368
Deflection, d, mm

Figure 3. Typical loading behavior at 4 K, showing serrations.

5.5 Owing to adiabatic heating, the fracture resistance of materials at

4 K may be affected by rapid loading. For this reason, 4-K tests are performed at low loading rates (see section 9.4). The possible effects of

heating must be considered when applying standard measurements to the design

of practical structures subject to high loading rates.

5.6 The parameters Kj c , J c , and Jj c are suitable for characterizing

a wide range of materials at 4 K. Nevertheless, some materials cannot be

tested satisfactorily by the methods of this standard. Some materials are too brittle to be properly fatigue -cracked, or too thin (B < 1.6 mm) for Kj c testing (see Methods E 399-83, section 1.1). Others are too tough for valid

Jj c tests (see Methods E 813-81, section 4.3). Retesting with a change in

specimen dimensions may provide valid data in some cases, but there is no assurance that valid results for a given material will be obtained by the

procedures of this standard.

369
6 . Apparatus

6.1 Test Machines—The tests shall be performed using screw-driven or servohydraulic tensile test machines that meet the requirements of Practices

E 4 regarding the verification of load accuracy. The equipment must include

a means for autographic recording of the applied loads and specimen crack

opening displacements. The compliance of the test machine and its cryostat

should be known; it can be measured at 4 K by replacing the specimen with a rigid block.

6.2 Cryogenic System Design—For conventional test machines having a maximum load capacity of 10 tons, it is recommended that the cryogenic apparatus be designed to accommodate a standard, 25 -mm- thick compact specimen

(see section 8.1.3). A servohydraulic test machine with an inverted stage

(actuator mounted above the specimen) and a cryogenic loading frame that is

free at the lower end is also recommended. Such a system is nearly optimum

in regard to operating convenience and cost effectiveness for testing a wide

range of materials at 4 K.

6.3 Cryogenic Construction Materials—Many common structural materials,

including most steels, are brittle at 4 K. Therefore cryostat components

must be well designed and fabricated using cryogenic materials of suitable

strength and toughness. Austenitic stainless steels are generally recom- mended. Maraging steels, nickel plated to prevent rusting, are useful for

specimen clevises. Alpha-titanium alloys are appropriate for applications

requiring high strength- to- thermal conductivity ratios, such as pull rods.

6.4 Dimension- Measuring Devices—Micrometers and other devices used

for measuring the dimensions of specimens shall be accurate and precise to

at leash one-half the smallest unit to which the individual dimension is

required to be measured.

6.5 Cryostats and Support Apparatus

6.5.1. Cryostats-—A cryostat capable of retaining liquid helium for the

duration of testing is required. Usually, cryostats are custom-built, but

they may be designed for use with commercially available Dewars. Efficient

designs involve thermal and mechanical trade-offs (14). Several designs are described in the literature, including turret-disc devices for testing many

specimens successively with one cooling (7,12,14—16).

6.5.2 Dewars—Stainless steel Dewars are recommended for low maintenance, durability, and safety. Glass Dewars are fragile and hazardous, and

fiberglass Dewars are relatively expensive. A single Dewar is generally

sufficient for tests of short duration, but double-Dewar arrangements using

an outer Dewar of liquid nitrogen to surround the inner Dewar containing

liquid helium are also possible (14—16).

6.5.3 Ancillary Equipment—-The Dewars and transfer lines for handling

liquid helium must be vacuum insulated. Pressurized gases, vacuum pumps,

and liquid nitrogen apparatus are therefore required. After testing, the

370
helium may be released as a vapor to the atmosphere, recycled, or reliquefied. Recycling or reliquefaction requires substantial investment in purification and cryogenic systems.

6.6 Clevis Grips and Pins:

6.6.1 Design—Conventional clevis grips conforming to the design given

in Fig. 3 of Methods E 813-81 may be used. The loading pin holes have flats

flats to allow specimen rotation during testing. To minimize friction, the

flats must remain free of plastic indentations. Alloys with adequate 4-K

mechanical properties are required for construction.

6.6.2 Construction—Materials having adequate 4-K strength, toughness,

and fatigue resistance are required for grips and pins. The maraging steels

(i.e., 250 and 300 grades) can be used for grips of the conventional design

(7,14—16). If lower strength materials such as AISI 304LN are used, the

thickness of the clevis arms should be enlarged. Pins must be manufactured

from high-strength materials to resist bending.

6.7 Extensometer:

6.7.1 Requirements—An extensometer device is required to measure edge

and load-line displacements for Kj c or J tests, respectively. In addition to

rendering accurate displacement measurements at 4 K, the device must be able

to sustain cycles of thermal shock, moisture condensation, and mechanical

fatigue

.

6.7.2 Types—Conventional clip gages, such as those described in Methods E 399 and E 813, may be purchased or built using cryogenic materials and

fabrication techniques. Ring-shaped clip gages of various sizes are recom- mended for cryogenic testing (17) . The gage design recommended in Methods E 399 may also be used if the total displacement is expected to be 2.5 mm or

less. For a larger linear working range, enlarged gages are necessary (see

Fig. 4 of Methods E 813).

6.7.3 Construction—A satisfactory combination of active strain gage

elements, backing materials, and bonding agents must be used for cryogenic

extensometers . One common choice is a Ni—Cr-alloy gage with a temperaturecompensated active element. Typically, 120- or 350-0 gages are wired in a Wheatstone bridge arrangement. Electrical wires insulated with polytetraf luoroethylene and soft solders with low tin contents are recommended to

reduce mechanical failures and electrical shorts at 4 K.

6.7.4 Calibration—The clip gage or other extensometer device must be

calibrated at room temperature and at 4 K. This may be done with a barrel

micrometer device fitted with vertical extension tubes with the gage mounted

at the lower end and immersed in liquid helium for calibration at 4 K (17).

Once the calibration is known and proven linear, accurate, and reproducible,

room- temperature checks may be performed to verify the 4-K calibration

indirectly. Periodically, the 4-K calibration must be directly verified,

especially when damage is suspected or repair has been performed.

371
6.8 Temperature Maintenance and Liquid-Level Indicators—The intended

cryogenic environment is maintained by the fixed boiling point of the liquid

helium bath. A thermocouple on the specimen is not necessary, but a suitable

meter or indicator is required to ensure the proper liquid level. An on—off

carbon-resistor meter, a continuous- level meter with superconducting wires,

or a cryogenic thermocouple (Chromel-vs-gold + 0.07% iron is most suitable

below 20 K) shall be positioned in the cryostat above the specimen to ensure

that it remains fully submerged for the duration of the test.

7 . Sampling

7.1 Material Condition—Samples for specimen preparation shall be taken

from the material in its final metallurgical condition. All heat treatments

shall be performed prior to sampling and specimen machining.

7.2 Sample Location—The sample location shall be chosen to be representative of the stock material. Allowance must be made for areas affected

by the cutting process.

7.3 Specimen Orientation—The orientation of the specimen notch plane

and the direction of load application relative to the principal stock working directions must be stated using the two- letter code of Methods E 399-83,

section 5.1.3. The TL orientation is favored for general material evaluations and alloy development at 4 K.

8. Specimens

8.1 Specimen Configuration:

8.1.1 Specimen Types—Any specimen configuration cited in Methods E 399-

83 or 813-81 may be used at 4 K. Current practice, however, favors compact

and bend specimens. In comparison, bend specimens require greater mass for an equivalent ligament size. In addition, the long span of the bend specimen

transverse to the load axis requires fixtures that are inconvenient for some cryostats. Compact specimens, therefore, are usually preferred at 4 K.

8.1.2 Standard 4 -K Test Specimen—The standard compact specimen for

material evaluations at 4 K has a thickness, B, of 25.4 mm and a width, W, of

50.8 mm. The W/B ratio is 2.0, and the planar dimensions are proportional to

W. This specimen is large enough to provide valid data for most austenitic

stainless steels tested at 4 K. The specimen may be modified for edge or load-line displacement measurements, as shown in Fig. 4.

8.1.3 Alternative Specimens—If the standard specimen of section 8.1.2

is not practical for any reason, an alternative geometry or size may be

selected, following Methods E 399 and E 813. Compact specimens with W/B - 2 and planar dimensions proportional to W are generally recommended, but the

choice of specimen type and size should be based on the technical purpose,

the load limit of the test machine, the stock dimensions, and the need to

obtain representative samples of the material being tested.

372
63.5 Dimensions

in mm

Figure 4. Alternative compact specimens of standard size; see also

Methods E 399-83 and E 813-81 for acceptable specimen geometries

373
8.2 Specimen Preparation :

8.2.1 Starter Notch—The notch should be machined according to the

guidelines of Methods E 813-81. The notch should be centered along the

midplane of the specimen to within 1 mm. The angle between the notch plane

and the midplane of the specimen should not exceed 1 degree.

8.2.2 Load-line Notch Modification—For J tests, a load-line notch

modification (see Methods E 813-81) for clip-gage attachment between razor blades is recommended. Integrally machined knife edges are also possible,

but they may tend to increase friction effects. If the displacement is not

directly measured at the load-line in Jj c tests, then the method of determining the displacement at the load-line from measurements at the edge location must be described.

NOTE 2—If a load-line modification is used and a Type I load—displacement curve is obtained, then Kp c may be calculated as usual.

8.3 Fatigue Precracking :

8.3.1 General Requirements—All specimens must be carefully fatiguecracked. During fatigue, the following requirements always apply:

( 1 )

( 2 )

e3)

The maximum stress-intensity factor, Kmax , during fatigue crack

extension shall not exceed 33 MPa-m*

The load range (AP) during fatigue shall not be less than 90% of

the maximum fatigue load, Pmax . Pmax must not exceed 0.4Pl, where Pl is the limit load for the

material at the fatigue test temperature. For compact specimens:

PL - Bb2oy/(2W + a)

8. 3. 1.1 High- toughness steels. If the 4-K fracture toughness exceeds

Kj c - 100 MPa-mn

, the fatigue cracking may be conducted at room (295 K) or cryogenic temperatures (77 or 4 K) , and the conditions of section 8.3.1 must

be satisfied.

8. 3. 1.2 Low- toughness steels. If the 4-K fracture toughness is Kj c -

100 MPa-m^ or less, the recommended temperature for fatigue cracking is

4 K, and the conditions of section 8.3.1 apply. Also, the maximum stressintensity factor in fatigue at 4 K, Kmax , must not exceed 60% of the material's Kq c value at 4 K. If the fatigue cracking is performed at a temperature higher than 4 K (i.e., 295 K)

,

then it must be demonstrated that the

procedure has no effect on subsequent toughness measurements at 4 K.

NOTE 3 —Sizable reductions of the maximum load during the

fatigue precracking should be avoided since they may retard

the crack growth rate or increase the degree of crack- front

curvature

.

374
8.3.3 Fatigue Crack Length—The ratio of the initial fatigue crack

length to specimen width, a/W, shall be in the range from 0.45 to 0.70. The

optimum value is a function of the material behavior. If a Type I test

record is expected and Kj c can be measured directly, then a/W should be 0.45

to 0.55, compatible with Methods E 399-83. If nonlinear behavior is

expected and a resistance curve can be determined, the crack- length ratio

should be 0.6. For exceptionally strong and tough steels, a ratio near 0.7

may be necessary to meet the test-machine load limitations. For the initial

test of an alloy of unknown toughness, a ratio of 0.6 is recommended; this can be adjusted in retests, if necessary.

8.4 Side Grooves :

8.4.1 Acceptable Uses Side grooves shall not be used for materials

that display Type I or Type II test records. For Type III test records,

side grooves may be used to alleviate fatigue crack- front curvature and

crack- front tunneling, or to improve the agreement between the optically measured and compliance-calculated end-of-test Aa values.

8.4.2 Machining—Side grooves, if used, shall be machined after pre- cracking, following Methods E 813-81, section 7.5.1. For standard compact

specimens, which have a thickness of 25 mm, the recommended side-groove

geometry is shown in Fig. 5. For standard specimens, the maximum depth of

each groove shall be 2.5 mm (a total depth of 20% of the gross specimen

thickness)

.

A B C

Figure 5. Recommended side-groove geometry: A) plane view, B) side view.

and C) details for a standard 25-mm-thick compact specimen.

375
9. General Procedures

9.1 Specimen Measurement—Measure the specimen dimensions according to

Methods E 399-83 (section 8.3) and E 813-81 (section 8.8).

9.2 Installation—Install the specimen in the apparatus, leaving slack

for the instrumentation wires so that they will be neither stretched nor crimped during positioning of the Dewar and testing. Before cooling, check

the seating of the specimen and grips at room temperature. The specimen may

be preloaded slightly for this purpose.

9.3 Cooling—First, dry the apparatus thoroughly with pressurized air

or a heating gun. Also, purge any moisture from the cryogenic transfer

lines. Ice can block the transfer lines or cause erratic loading behavior

if it forms on the loading pins, clevises, specimens, clip gage, or loadtrain members. Next, position the Dewar and precool the apparatus to 77 K

by transferring liquid nitrogen into the cryostat. After reaching equilibrium at 77 K, remove all the liquid nitrogen and transfer liquid helium

into the cryostat. Add sufficient liquid helium to compensate for evaporation during the time required to complete the test. Before testing, hold

each specimen of standard size in the liquid helium bath for a minimum of

15 min.

NOTE 4—The heat- transfer characteristics of gaseous

helium are inferior to those of liquid helium; therefore

full submersion in the liquid is required to minimize

specimen heating.

9.4

Loading :

9.4.1 Control Mode—For tests with a screw- driven machine, the rate of

crosshead movement shall be controlled. For tests with a servohydraulic

machine, either stroke control or clip-gage displacement control may be

used; load control, however, should not be used.

9.4.2 Load Rate—Load the specimen at a rate such that the time, tp, to

reach the load at first deviation from the linear portion of the test record

is between 0.5 and 10 min. For standard 25 -mm- thick compact specimens, the

crosshead rate should be 0.5 mm-min-1 or less. For additional details of

test procedure, see Methods E 813-81, section 8.

9.5

Test Record—Record the load- cell and clip -gage outputs autographicaliy, following the guidelines of Methods E 399-83, section 8.4 for Kj c tests, or as appropriate for Jj c tests. For Jj c tests, a recording of the

computer -calculated J—Aa curve in real time is also recommended.

376
10 Interpretation and Calculation of Result:

c ..a - _ Cc C

10.1 General—Classify the test record as one of th

shown in Fig. 2. Depending on the load-vs -displacement

which fracture criterion is to be used. More than one fracture tar a: may be applied in some cases.

10.2 Kjc measurement—For a Type I test record, calculate the f: (Kq or KIc ) according to Methods E* 399. Construct the secant offset"

calculate the conditional toughness value, Kq, and determine if it c. as a valid Kjc according to Methods E 399-83, section 9. If the K

*- fails to qualify as a valid KIc measurement, calculate J c accordin'

section 10.3 of this standard.

NOTE 5—Notch modifications for load-line

measurements may invalidate K^ c calculate

may be necessary retest some specimens wit

modifications to prove the validity of K-.

— — a - d w c — tr . . .

- U.-S d - 4 K. It

— SdS’j.r’szLsr1. “ s

10.3 Jc measurement—For Type II tests producine nonlinear load-vs- displacement curves, determine J c as follows: Calculate the value of she J- integral corresponding to the total area under the load-vs -displaceme-

curve up to the critical pop-in or fast fracture event The area o*‘ "^ter est is cross-hatched in Fig. 2. The measurement shall be accurate :0 it may be performed using a polar planimeter or numerical integration, o^ --I

load-vs -displacement signals by computer techniques. For comrac -

s-

use the following expression to calculate the value of J :

- 6 . .m>

where

•J c “ - ^o/T|11 )

‘ A/3b

a0 W

f(a 0 /W)

B

b

original crack length, including precrack len.tth.

specimen width,

a dimensionless coefficient that corrects for the tensile

component or loading, as given in section 9

* 1 a~d ~

of Methods E 813-81, area under the load-vs -load- line-displacement record

in energy units

,

specimen thickness, and

initial uncracked ligament, V - a-

.

10.4 Jjc measurement—For Tvoe III r:e^= - ... _ tance curve technique, where J is plotted versus physical crack growth ^ • rne single-specimen unloading compliance method is recommended as -_-e emcient than the multiple-specimen technique at 4 :< because it reduc-T- required number of test specimens and the cryogenic fluid costs. Quid, ror the comp.iance test method are given in the arcer.dix of '•let'-'d

“

~ and m references (2-7).

377

(

V

l>0
10.4.1 Using the expression J = f(a/W)’A/Bb, calculate J at points

along the load-vs - load- line—displacement curve as a function of the crack

length inferred from compliance. For side-grooved specimens use net thickness, Bn , instead of B for J calculations and use the effective thickness,

Be ff, instead of B for crack growth predictions from unloading compliance

data (see Methods E 813-81 sections 8.7.1 and 9.1.1).

10.4.2 Prepare a plot of J-vs-Aa

p spanning the 0.15—1.5 -mm exclusion

interval, as given by Methods E 813-81. Calculate the area under the load- vs -displacement curve, including the area under serrations. For tests in

clip -gage displacement control, the area under a serration is zero. Thus,

tests in stroke control are expected to give higher values of J for a given

crack extension, owing to the area increase associated with serrations. The

intent is not necessarily to measure a physically meaningful J—Aa curve, but

rather to obtain a regular data set, enabling linear extrapolation to the

initiation of crack extension at djc10.4.3 Construct the blunting line (assuming J * 2c?yAa) , plot the four

point linear regression line, and qualify the J—Aa calculations following

Methods E 813-81, section 9.2 in every respect.

NOTE 6—Compliance data usually show an anomalous

negative crack growth effect in early stages of loading

Correct for this effect by shifting the blunting line

along the Aa axis until it coincides with the upper

portion of J—Aa data, which indicate a normal bluntingline trend. The negative crack growth data are thus

ignored (see Fig. 6).

10.4.4 Determine Jq as the value of J at the blunting- and regressionline intersection. Qualify the data and determine whether the Jq is a valid

Jj c determination according to Methods E 813-81, sections 9.3 and Al

.

Figure 6. Method of correction for negative crack growth.

378
11. Report

11.1 Material Characterization:

11.1.1 Material Description—Describe the material tested. Include the

product form (e.g., plate, forging, casting), processing, heat treat-ment,

and any other relevant metallurgical information.

11.1.2 Tensile Properties Report the yield strength, tensile strength,

and elastic modulus for the tested material at room temperature and at 4 K.

11.2 Specimen Characterization:

11.2.1

Report the crack-plane orientation using the identification code of Methods E 399-83.

11.2.2

Report the specimen dimensions, including thickness, width, initial ligament size, and initial crack length.

11

; 2 - 3 For side-grooved specimens, report the side-groove geometry, including the included angle, notch- tip radius, and net section thickness.

11.3 Fatigue-Cracking Characterization:

11.3.1

Report the maximum fatigue load, the maximum stress - intensity factor, the stress-intensity-factor range, and the temperature used for the fatigue precracking.

. 11,3,2 Report the final fatigue crack length obtained using the three- point average for KIc tests or the nine-point average for J Tr tests as de- scribed in Methods E 399-83 or E 813-81, respectively. Also'meLure the surface crack lengths, and report the maximum deviation of any single measurement from the average.

TeS

l f

ecord—IAscribe Che observed loed-vs -displacement behavior, classifying it by type according to the examples shown in Fig. 2.

11.5 Loading Rate—Report the loading rate in terms of tD or the rate

ine Soecifv

stress -^tensity factor during the region If linear loadwaf’usfd

^ er crosshead . stroke, or clip-gage displacement control

11.6 Fracture Toughness Measurements:

11.6.1

Report the compliance of the test machine and cryostat used.

Kt

3

^ 1 t6St record ’ reP° rt the toughness in terms of Kq or u e an assessment of the validity according to Methods E 399-83

cen J

1,6

; 3

/°f. Type 11 test records, report the value of J c as defined in on o this standard. If a compliance technique was used for J-Aa

379
measurements, estimate and report the amount of subcritical crack extension

(if any) that occurred prior to the point of instability.

11.6.4 For Type III test records where a full resistance curve can be

developed, report the value of Jj c and calculate Kj c (J). Assess the validity of Jj c according to the criteria of Methods E 813-81.

11.7 Fracture Surface Appearance—Describe the fracture -surface appearance. For computerized Jj c tests, report the end-of-test crack extension

increment. Also report the averaged values of the compliance -predicted and

the optically measured final crack lengths.

11.8 Replicate Tests—State the number of replicate tests performed,

the average fracture toughness values, and a measure of scatter.

11.9 Anomalies—Report any anomalies observed in the material behavior,

test records, or failure mechanisms.

12. Accuracy and Precision

12.1 Accuracy—There is no accepted standard fracture toughness value

for any material. In the absence of such a true value, any statements about

accuracy are meaningless.

12.2 Precision—Information on the precision of 4-K fracture tests is

being obtained from interlaboratory test programs currently in progress.

REFERENCES

(1) Cryogenics Safety Manual—A Guide to Good Practice , 2nd Ed. , Mechanical Engineering Publications, London, 1982.

(2) Clarke, G. A., Andrews, W. R. , Paris, P. C., and Schmidt, D. W.

,

"Single Specimen Tests for Jj c Determination," in: Mechanics of

Crack Growth, ASTM STP 590, American Society for Testing and

Materials, Philadelphia, 1976, pp. 27—42.

(3) Joyce, J. A. and Gudas , J. P., "Computer Interactive Jj c Testing of

Navy Alloys," in: Elastic-Plastic Fracture, ASTM STP 668, J. D.

Landes, J. A. Begley, and G. A. Clarke, eds., American Society for

Testing and Materials, Philadelphia, 1976, pp . 451—468.

(4) Clarke, G. A., "Single Specimen Tests for Jj c Determination

Revisited," in: Fracture Mechanics: Thirteenth Conference, ASTM STP

743, R. Roberts, ed. , American Society for Testing and Materials,

Philadelphia, 1983, pp . 553-575.

380
( 5 ) Tobler , R. L. , "Ductile Fracture with Serrations in AISI 310S 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., American Society for Testing and Materials, Philadelphia 1983

pp. 763-776.

(6) Shimamoto, S., Nakajima, H. , Yoshida, K. , and Tada, E. , "Requirements

for Cryogenic Structural Alloys for Superconducting Magnet Cases,"

in: Advances in Cryogenic Engineering—Materials, Vol. 32, Plenum

Press, New York, 1986, pp . 23-32.

(7) Shimada, M. , Ogawa, R. , Moriyama, T. , andHoriuchi, T. , "Development

of a Cryogenic Fracture Toughness Test System," Cryogenic Technology

Vol. 21, No. 5, 1986, pp . 269-274.

(8) Takeuchi, M. , Shoj i , T. , Takahashi, H. , and Anayama, T. , "Evaluation

of fracture Toughness of Cryogenic Materials," Reports of the Research Institute for Strength and Fracture of Materials

,

Vol. 17

1984, pp. 1-12.

(9) Nakajima, H. , Yoshida, K. , Shimamoto, S., Tobler, R. L. , Purtscher,

P. T., and Reed, R. P. , "Round Robin Tensile and Fracture Test

Results for an Fe-22Mn-13Cr—5Ni Austenitic Stainless Steel at 4 K,"

in: Advances in Cryogenic Engineering—Materials, Vol. 34, Plenum

Press, New York, 1989, pp. 241-249.

(10) Ohnishi , K. , Shoji, T., and Takahashi, H. , "Effects of Loading Rate

and Machine Compliance on Heat Generation During Cryogenic Fracture

Toughness Testing," in: Workshop on Standardization of Fracture

Toughness Testing of Low Temperature Structural Materials, Tohoku

University, Sendai, Japan, 1986, pp. 27-35.

(11) Shimada, M. , Tobler, R. L. , Shoji, T. , and Takahashi, H. , "Size, Side -Grooving, and Precracking Effects on J Ic Data for an SUS 304

Stainless Steel at 4 K," in: Advances in Cryogenic Engineering— Materials

,

Vol. 34, Plenum Press, New York, 1988, pp . 251-258.

(12) Ogata, T., Ishikawa, K. , Yuri, T. , Tobler, R. L. , Purtscher, P. T.

,

Reed, R. P. , Shoji, T., Nakano , K. , and Takahashi, H. , "Effects of' Specimen Size, Side-Grooving, and Precracking Temperature on J- Integral Test Results for AISI 316LN at 4 K," in: Advances in Cryogenic Engineering—Materials

,

Vol. 34, Plenum Press New York

1988, pp. 259-266.

U3) Yoshida, K. , Nakajima, H. , Oshikiri, M. , Tobler, R. L. , Shimamoto, S .. ,^ra ’ R ;

a?

d Ishizaka, J., "Mechanical Tests of Large Specimens at

. Facilities and Results," in: Advances in Cryogenic Engineer inr— Materials, Vol. 34, Plenum Press, New York, 1988, pp . 225-232.

381
( 14 ) Fowlkes, C. F. and Tobler, R. L. , "Fracture Testing and Results for

a Ti—6A1—4V Alloy at Liquid Helium Temperature," Engineering Fracture

Mechanics

,

Vol. 8, 1976, pp. 487—500.

(15) 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.

(16) Juhas, M. C., Duncan, D. B. , and Streit, R. D. , "Development of a Facility for Fatigue and Fracture Testing of Stainless Steels at

Liquid Helium Temperature," Cryogenics , Vol. 23, 1983, pp . 603-607.

(17) Tobler, R. L. and Shepic, J. A., "Design and Performance of a Ringshaped Clip Gage for Fracture Mechanics Testing," Journal of Testing

and Evaluation

,

Vol. 13, 1985, pp . 299—302.

382
JAPAN—U.S. COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

FIFTH DRAFT

SEPTEMBER 1987

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 sjimiiar to that of other ASTM tensile test standards, but the contents in- clude modifications for cryogenic testing, which requires special apparatus smaller specimen size, and concern for serrated yielding, adiabatic heating’ and strain-rate effects.

To conduct a tensile test by this standard, the specimen in a cryo- stat 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 tensUe strength elongation, and reduction of area. The determination of the elastic modulus is treated in Method E 111.

It o 1,4 Values s 5

ated 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

TMriLeteStfn?' IC " ^

USer ' S to learn and obHTM

end and

f 7 Prac tlces . Safety guidelines for handling liquid helium 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 eievation; it is 4.0 K at NBS in Boulder, Colorado, 1677 m (5500 ft) above sea level. In this standard, the test temperature is nominally designated 4 K.

383
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.

384
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 plas- tic 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

t

FIG 1 Example of a Cryostat with Double -Dewar Arrangement for Tensile Testing at 4 K

385
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

.

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 that of the cryogen at 4 K

during the discontinuous yielding events (see Fig. 2). This behavior varies

with specimen size and testing speed, but altering the mechanical test variables cannot eliminate the periods of internal heating (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.

386
STRESS,

MPa

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 de- sign in applications when actual circumstances may approach load-controlled conditions

.

FIG 2 Engineering Stress-Strain Curves for an AISI 304L Austenitic Stainless Steel at 4 K

387

TEMPERATURE,
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—6A1—4V

and Ti—5A1—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

.

388
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. 2 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 mea- suring 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 he 1 inn

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

.

389
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.

Red (+ excitation)

n (constant voltage)

Room Temperature

Dummy Resistor *

Green

(+ signal)

Room Temperature

Dummy Resistor*

Shunt Resistor ( excitation)

FIG 3 Method of Strain-Gage Wiring for Cryogenic Tests

390
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.

391
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 re- quirements 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 are 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 con- ventional 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.

7. 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.

392
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)

393
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.

394
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 con- trolling 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 one- half 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.

395
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 re- sult, yield strength data from load—time curves may be less

accurate than those of the recommended method.

8.6.3 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.

8.6.4 Elongation—Calculate the percentage increase of elongation

according to Methods E 8 or E 8M, section 7.6.

8.6.5 Reduction of Area—Calculate the percentage of reduction of area according to Methods E 8 or E 8M, section 7.7.

8.6.6 Rounding Reported Test Data—-Round-off the calculated numerical

test results according to Methods E 8 or E 8M, section 7.8.

8.6.7 Replacement Specimens—If necessary, discard any invalid data and

test replacement specimens as per Methods E 8 or E 8M, section 7.9.

396
9 . Report

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 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.

10.2 Bias—The bias of these methods includes quantitative estimates

of uncertainties of the dimension-measuring devices, the calibration of rest

equipment, and the skill of the operators. At present, bias statements

should be limited to the documented performance of individual laboratori*.

397
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.

(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.

398
U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

1. PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No. 3. Publication Date

SHEET (See instructions) NBSIR 88-3082 May 1988

4. TITLE AND SUBTITLE

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures — XI

5. AUTHOR(S)

R. P. Reed, Editor

6. PERFORMING ORGANIZATION (If joint or other than N BS, see instructions) 7. Contract/Grant No. national bureau of standards

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

9.

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

Department of Energy

Office of Fusion Energy

Washington, DC 20545

10. SUPPLEMENTARY NOTES

Previous reports in this series: 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 computer program; SF-185, FIPS Software Summary, is attached.

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

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 magnetic 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 Bureau of Standards and the Office of

Fusion Energy of the Department of Energy; it is managed by NBS and sponsored by

DOE. Research is conducted at NBS and at other laboratories through subcontracts with NBS.

Research results for 1987 are presented in technical papers under five

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

Nonmetallics , 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 (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words t y -.r~ . austenitic steels; composites; cryogenic properties; elastic properties; low temperatures; mechanical properties; nonmetallics; stainless steels; structural alloys; welds

13.

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