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<c'. V; MATERIALS STUDIES FOR MAGNETIC

FUSION ENERGY APPLICATIONS

AT LOW TEMPERATURES - XIII

TEMPIHAT0RI

NISTIR 3944

MATERIALS STUDIES FOR MAGNETIC

FUSION ENERGY APPLICATIONS

AT LOW TEMPERATURES - XIII

Edited by

R.L. Tobler and R.P. Reed

National Institute of Standards and Technology

Boulder, Colorado 80303-3328

September 1990

Prepared for

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

U.S. DEPARTMENT OF COMMERCE. Robert A. Mosbacher. Secretary

NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY. John W. Lyons. Director

CONTENTS

SUMMARY vi

ORGANIZATIONAL CONTACTS vii

PROGRAM DESCRIPTION viii

STRUCTURAL ALLOYS

Structural Alloys Program 2

Mechanical Properties of a Superconductor Conduit

Proposed for ITER Magnets 7

R.L. Tobler, R.P. Walsh, C.N. McCowan, and R.P. Reed

Near-Threshold Fatigue Crack Growth Rate Measurements for Five

Austenitic Stainless Steels at 4 K 27

R.L. Tobler

Creep of Copper: 4-295 K 51

R.P. Reed, N.J. Simon, and R.P. Walsh

Strain Rate Effect on Tensile Properties at 4 K of a VAMAS

Round-Robin Austenitic Steel 61

R.P. Reed, R.P. Walsh, and R.L. Tobler

Load- Controlled Tensile Tests of Austenitic Steels at 4 K 69

H.M. Lee, R.P. Reed, and J.K. Han

Effect of Processing on the 4-K Mechanical Properties of a

Micro-alloyed Austenitic Stainless Steel 79

P.T. Purtscher, M.C. Mataya, L.M. Ma, and R.P. Reed

Austenitic Steels with 9% Cr for Structures at Cryogenic

Temperatures 87

P.T. Purtscher, M. Austin, C.N. McCowan, R.P. Reed,

R.P. Walsh, and J. Dunning

Micromechanisms of Ductile Fracture and Fracture Toughness

in a High Strength Austenitic Steel 95

P.T. Purtscher

iii
WELDING

Welding Program 219

Development of an SMA Electrode to Match Type 316LN

Base Metal Cryogenic Mechanical Properties 223

T.A. Siewert and C.N. McCowan

Cryogenic Mechanical Property Data for Gas Metal Arc Welds

with a Composition of 20Cr-25Ni-4. 5Mo 233

T.A. Siewert and C.N. McCowan

Fracture Characterization of 316LN Primary Austenitic GMA Welds . . 239

C.N. McCowan and T.A. Siewert

TEST STANDARDS

Development of Test Standards Program 251

VAMAS Technical Working Party Business Meeting 253

R.L. Tobler

The Effect of Side Grooves on Kj c Test Results

for High Strength Austenitic Stainless Steel at 4 K 255

R.L. Tobler

Interlaboratory Tension and Fraction Toughness Test Results for

CSUS-JN1 (Fe-25Cr- 15Ni-0 . 35N) Austenitic Stainless Steel at 4 K . . 261

H. Nakajima, K. Yoshida, S. Shimamoto

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

Jj c Data Analysis Methods with a "Negative Crack Growth"

Correction Procedure 269

Y.A. Rosenthal, R.L. Tobler, and P.T. Purtscher

Charpy Impact Tests of Superconductor Conduit Alloy Welds

Near Absolute Zero 273

R.L. Tobler and D.P. Vigliotti

Comparisons Between Jj c and Charpy Absorbed Energy as Measures

of Cryogenic Fracture Toughness 283

I.S. Hwang, M.M. Morra, R.G. Ballinger,

H. Nakajima, S. Shimamoto, and R.L. Tobler

Tension Testing of Structural Alloys in Liquid Helium 301

(ninth draft, ASTM Subcommittee E 28.10)

Proposed Revision of ASTM Standard Designation E 23-88,

"Standard Test Method for Notched Bar Impact Testing of

Metallic Materials," 321

(ASTM Task Group E 28.07.02)

iv
Proposed Annex A3 to ASTM Method E 813-88, "Test Procedures

for Austenitic Alloys in Liquid Helium at 4 K," 325

(ASTM Subcommittee E 24.08.04)

TECHNOLOGY TRANSFER

Technology Transfer Program 341

Twelfth Cryogenic Structural Materials Workshop 343

N . J . S imon

International Thermonuclear Experimental Reactor

Specialists Meeting on Magnet Materials 355

R.P. Reed, N.J. Simon, and R.L. Tobler

APPENDIX

Mechanical Properties of Ultra-Low Carbon Steel and Effects

of Warm Prestress at 4 K 363

R.L. Tobler, L.M. Ma, and R.P. Reed

Cryogenic Fatigue of High-Strength Aluminum Alloys

and Correlations with Tensile Properties 375

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

v
'
SUMMARY

This report contains the results of research 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. This

research facilitates their design and development. The program was

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

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

is managed by NIST and sponsored by DOE. The research is conducted at NIST

and at other laboratories through cooperative agreements with NIST. Some

research, such as the development of test standards, is partially funded by

NIST through its own base research programs. Cooperative research at other

laboratories is performed by their own funding.

Research results for 1989/90 are presented in technical papers under four

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

Test Standards and Technology Transfer. Objectives and research highlights

are summarized in the introduction to each program area.

This report also includes an Appendix consisting of two papers with data for

a ferritic steel and four aluminum alloys. That research was funded

primarily by other programs but the information is of interest to magnetic

fusion energy applications.

NOTE : Certain commercial equipment, instruments, or materials are

identified in this report to specify the experimental procedure adequately.

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

specific research. In no case does such identification imply recommendation

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

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

best available for the purpose.

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 of 1990.

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

Program Monitor M. Cohen (301) 353-4952

National Institute of Standards

Program Manager

Structural Alloys

Welding

Technology Transfer

Test Standards

and Technology, Boulder, CO 80303

R.L. Tobler (303) 497-3421

R.P. Reed (to Feb 90) (303) 497-3870

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

C . N . McCowan (303) 497-3699

T.A. Siewert (303) 497-3523

E.S. Drexler (303) 497-5350

N.J. Simon (303) 497-3687

R.P. Reed (303) 497-3870

R.L. Tobler (303) 497-3421

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 lowtemperature design, construction, and measurement programs, such as the

Large Coil Project (LCP) , the Compact Ignition Tokamak (CIT) , and the

International Thermonuclear Experimental Reactor (ITER), we are aware of new

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

research on stainless steels in support of ITER designs.

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
m
STRUCTURAL ALLOYS

STRUCTURAL ALLOTS PROGRAM

LEADER: R.P. Reed

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

OBJECTIVES

:

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

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

pulsed superconducting magnets

o Development of effective codes and standards for low- temperature

property measurements and structural design

RESEARCH HIGHLIGHTS:

o Selective mechanical property tests of specific candidate alloys for

magnet applications were carried out. These include tension, fracture

toughness, and fatigue crack growth tests of JK1 (modified 316LN alloy)

for superconducting conduit; fatigue crack growth and threshold fatigue

tests for AISI 316LN and related alloys for magnet cases; and creep

tests of copper for conductor and stabilizer applications.

o A computer-controlled test apparatus was used to measure fatigue crack

growth rates and threshold stress intensity factors for austenitic

stainless steels at room temperature, 76, and 4 K. Long-crack

thresholds were determined using a decreasing AK technique and the

database at 4 K was augmented by tests of 6 alloys. A simulation

technique was also developed and used to characterize the behavior of

short cracks. Short-crack simulation tests of JK1 conduit steel are

reported and tests of other austenitic stainless steels are in progress

to develop a comprehensive database.

o Selected papers in this section describe various effects (strain rate,

load control, processing, composition, test temperature, and stress

state) on the mechanical behavior of stainless steels. These studies

impact cryogenic technology by improving our understanding of material

behavior; they contribute to a potential lowering of material

costs, and to resolution of technical points regarding the development

of cryogenic test standards and design codes.

3
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 K

Fatigue Crack

Growth Rate

295. 76. 4 K

Fatigue Report

S-N Volume*

4 K (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-13Cr-22Mn-0.2N

Fe-16Cr-8Ni-8Mn

Japan

U.S.

VI

VIII

XII

VIII

VII (65)

XII (239)

VIII (151,181)

Fe- 16Cr-28Ni P . R. China XII XII (141)

Fe-17Cr-18Mn-0.5N U.S. VII,

VIII

VIII VII (65)

VIII (181)

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

Fe- 1 7Cr-13Ni-2Mo

316

U.S. I, II,

III,IV

VI, VIII,

IX,X,XI

VI I, II,

VIII,

IX. X.XI

II, VIII,

IX

I I (15,71)

II (79)

III (49,105,117)

IV (147)

V (185)

VI (157)

VIII (181,209,251)

IX (15,27,43)

X (77)

XI (71)

Fe-17Cr-13Ni-2Mo

316LN

U.S. II, IX, XI II, IX, XI II II (79)

IX (15,27,53)

XI (71,103)

Fe- 18Cr-( 11-14) Ni-

(0-4) Mo

316LN

U.S. X X X X (9,31,43,55)

Fe-18Cr-3Ni-13Mn U.S. I,VI,

VIII, XI

III I, VIII, XI,

XII

I, VI, VII I (93)

III (91)

VI (53)

VII (85)

VIII (167,219)

XI (89)

XII (239)

Fe-18Cr-8Ni-

(0 . 6)Mn

U.S. X X (67)

Fe-19Cr-(6-15)Ni4Mn-(0 . 1-0 .3)

N

USSR VIII, XI VIII, IX, XI VIII (123)

IX (53)

XI (113)

Fe-19Cr-9Ni304 (cont'd

on next page

U.S. I, II,

III,

I, II III,I, II,

III,

II, III, IV I (15,71,213)

II (79,149,175)

III (15,91,105,117)

NBSIR 79-1609 (II); NBSIR 80-1625 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-(VI)

;

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

NISTIR 89-3931.

4
SUJ^AHY of MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

Country Fracture Fatigue Crack Fatigue Report

of Tensile Elastic Toughness Growth Rate S-N Volume*

Designation Supplier 295-4 K 295-4 K 295,76,4 K 295. 76, 4 R 4 K <J>age)

AUSTENITIC STEELS

Fe-19Cr-9Ni

304 , cont'd

U.S. IV, VI

VII,IX,X

IV, VIII IV, IX, X

IV IV (37,101,203,

215,227)

V (71)

VI (73)

VII (13,47,75,157)

IX (27) X (77)

Fe-19Cr-9Ni

304L

U.S. II, III,

IV, VIII,

IX,

X

i. II. IV II, IV,

VIII, IX

III, IV,

VIII

I, (213)

II (79,123,175)

IV (37,101,131,

215,302)

VIII (181,229,251)

IX (27,141)

X (9)

Fe-19Cr-9Ni-

(l-10)Mn

U.S. IV, V,

VIII

V IV, V, VIII,

IX

IV (77)

V (15,59,189)

VIII(181)

IX (53)

Fe-L9Cr-9Ni-N

304N

U.S. I. III.

IV, IX

IV I. III. IV, IX

I, III, IV I (93) III (15) IV (37,101,203,

215) IX (27)

Fe-19Cr-9Ni-N

304LN

U.S. II, III,

IV,V,VI

VIII, IX

IV, VIII II, III, IV,

VIII, IX

II, III, IV,

VIII

II (35,79) III (15) IV (37,101,203)

V (29)

VI (113)

VIII(181 ,207 ,229) IX (150)

Fe-19Cr-10Ni-

(0 .03-0. 29)N-

(0 .03-0 .09)C

U.S. II, VIII, VIII, IX,

XI, XII

III, VIII, III III (15)

VIII(145,181)

IX (159) XI (131)

XII (169,177)

Fe-( 14-19 )Cr-19Mn Switzerland XII XII (159)

Fe-20Cr-7Ni-8Mn

216

U.S. VIII VIII VIII(151 , 181)

Fe-20Cr-8Ni-2Mn

304HN

U.S. VIII VIII VIII(151 , 181)

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

VI (113)

Fe-21Cr-6Ni-9Mn U.S. I, VIII III, IV,

XII

I.VIII I I (15,71)

III (91)

VI (157)

VIII(151 ,181) IX (95) X (9)

XII (141)

Fe-21Cr-(6-12)Ni-

(cont.d on next

9Mn P.R. China

page)

XII XII (141,153)

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

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

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

NISTIR 89-3931.

5
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-21Cr-12Ni-5Mn U.S. I, II III, XII 1 ,11 1,11 I (93)

II (79) III (91)

XII (141)

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

310 VII, V (71,145)

VIII, VI (73)

IX,

X

VII (65)

Fe-19Ni-9Co U.S. IV IV (237)

200-300 grades

Reviews of IV,V,VI rv.v.vi IV.V.VI, IV, V.VI, IV, VI IV (17,257)

austenitic steels XI, XII XI, XII XI, XII VII V (171)

VI (11,127,157)

VII (103) XI (3,47)

XII (9,45,71,93)

ALUMINUM ALLOYS

Al-0.6Fe-0.2CuO.lSiU.S. IX IX (107)

1100

Al-2.7Cu-2.2Li u.s. XII X XII X (163)

2090 XII (131)

Al-6Cu-0 . 2Mn u.s. XII II II II II, XII II (19)

2219 XII (131)

Al-4Mg u.s. IV II II (35)

5083-0 IV (185)

Al-5.6Zn-2.5Mg- u.s. XII XII XII (131)

1 . 6Cu

7075

Al-5.7Zn-2.3Mg- u.s. XII XII XII (131)

1 . 5Cu

7475

COPPER ALLOY

99 . 99Cu-Ag u.s. IX, XI IX (95,107)

C10100, C10200 XI (147)

99 . 5Cu-Ag u.s. X X X X (83,127,139)

C10400

99 . 75Cu-Ag u.s. X X X X (83,127,139,

C15500 151)

97.6Cu-Ag-0.3CBe u.s. X X (83)

Cl 7510

99.95(Cu+Fe)-

18Cr-13Mn-3Ni0.37N

(laminate)

u.s. XI XI (159)

SUPERCONDUCTORS

Nb-45Ti u.s. III III (133)

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

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

;

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

NISTIR 89-3931.

6
MECHANICAL PROPERTIES OF A SUPERCONDUCTOR CONDUIT

PROPOSED FOR ITER MAGNETS

R.L. Tobler, R.P. Walsh, C.N. McCowan and R.P. Reed

Materials Reliability Division

National Institute of Standards and Technology

Boulder, CO 80303

A conduit for superconductor cables was tested to evaluate its

suitability for International Thermonuclear Experimental Reactor

magnets at liquid helium temperature. The conduit material, JK1

,

is a 316LN alloy modified by the addition of 0.05% Nb. Tension,

fatigue, and fracture toughness measurements are reported. A short

crack simulation technique was used to measure fatigue crack growth

rates and thresholds. At 4 K, the alloy's yield strength after

heat treatment exceeds the targeted 1200 MPa, but its toughness

(Jj c - 40 kJ/m2 for the T-L orientation) is lower than desired or

expected from alloy development studies. Metallurgical differences

between rolled plate and formed conduit, or between research grade

and commercial processing could account for the difference in

toughness

.

INTRODUCTION

High- field superconducting magnets use A15 type superconductors in

austenitic alloy conduit sheaths. The conduit sheaths protect the

superconductor cables and also contain the liquid helium coolant at 4 K. The

cable components and conduit are heated together during fabrication at

temperatures around 700 # C. This heat treatment reacts the Nb and Sn elements

to produce the superconducting compound NbjSn. It also ages the conduit and

adversely affects its mechanical properties.

The mechanical properties of austenitic stainless steels that are

proposed as conduit materials are degraded during aging by sensitization

(carbide precipitation) . Precipitation of M23 C 5 carbides on grain boundaries

is particularly deleterious. One method to combat this problem is to add

small amounts of Nb to retard the rate of carbide formation.

After testing laboratory plates, Shimada and Tone [1,2] decided that

alloy 316LN + 0.05% Nb offers better prospects than 304LN + 0.05% Nb or 22Mn

steel + 0.05% Nb. Despite some carbide and nitride precipitation during

aging, 316LN + 0.05% Nb retained a trans granular fracture mechanism at 4 K.

The alloy was designated JK1 and selected for further study. Recently, JK1

conduit was manufactured and proposed for International Thermonuclear

Experimental Reactor (ITER) magnets. In this paper, we report mechanical

property measurements for that conduit. Included are tension, fatigue, and

fracture toughness measurements at selected temperatures (295, 76, and 4 K)

before and after heat treatment.

7
MATERIAL AND SPECIMENS

Material was received in the form of a seamless rectangular tube with

the geometry shown in Fig. 1. In the as-received condition there is some

residual cold work from the extrusion process. As shown in Table 1, the

chemical composition is equivalent to plate compositions tested earlier.

Table 1. Chemical compositions of JK1 conduit (tested in this study) and JK1

research plate materials [1,2].

Material Cr Ni Mn C Mo N Nb Si S P

conduit 17.2 12.3 1.32 0.009 1.96 0.17 0.051 0.12 0.004 0.003

plate [ 1

]

17.17 12.14 1.48 0.010 2.03 0.19 0.055 0.10 0.0014 0.003

plate [ 2

]

17.12 12.45 1.26 0.010 1.98 0.198 0.049 0.12 0.005 0.007

The material was tested in the as -received (cold worked) and aged

conditions. To heat treat the material in laboratory furnaces, the 1 m long

stock was sectioned into 0.3 m lengths, heated in an argon atmosphere at

650°C for 180 h, and furnace cooled. This 650°C, 180 h treatment is

specified for the superconductor in ITER magnets.

Test specimens and dimensions are shown in Fig. 2. Flat specimens

(rectangular cross section) for tension tests were machined with their

loading axes in longitudinal or transverse orientations relative to the

extrusion direction or conduit length. Compact specimens for fracture

mechanics tests were machined in both the T-L and L-T orientations as defined

by ASTM Methods E 399-83. The specimen thicknesses were full size, that is,

identical to the thickness of the conduit wall; no machining was performed on

the specimen faces.

The microstructure of the conduit was examined before and after heat

treatment. Representative features after heat teatment are shown in Figs. 3

and 4. In Fig. 3, the grain boundaries are heavily decorated with

precipitates. Rod-shaped inclusions were observed in as-received material

and are also visible in the lower right portion of Fig. 3. Figure 4,

produced at lower magnification and with a different etchant, reveals banding

(alternate light and dark regions extending across many grains), which is

evidence of a deformed structure, chemical inhomogeneity, or both.

PROCEDURE

Tension tests were conducted using a screw- driven universal machine at a

nominal strain rate of 2.5 x 10" 4 s' 1

. The yield strength at 0.2% plastic

strain was measured using a 25.4 mm gage length, clip-on extensometer . The

cryogenic apparatus was described before [3].

8
55.2

Figure 1. JK1 conduit geometry and dimensions, as -received. The dimensions

are in mm, except for length which is m.

9
Note:

All dimensions in mm

Tolerance z 0.1 mm

Not to scale

55.8

44.7

11.1 —- 12.7 —

f

42.7

Figure 2. Test specimen geometries used in this study. The flat

tension (A) and compact specimen (B) thicknesses are 4.7 mm,

equal to the conduit wall thickness.

10
mE,

'3*'^ ^E^^Tz.. : Z* vTT^Sr*

^

*2 :.> 3

,.:r^2

-- -:• , -\

SSBH

\ . ,' .

— ' *r" -

Figure 3. Precipitates (presumably carbides) and inclusions in the JK1

conduit microstructure after aging at 650°C for 180 h

( lOOOx)

.

Figure 4. Banding in the JK1 conduit microstructure after aging at

650°C for 180 h (50x).

11
Fatigue and fracture toughness tests used a servohydraulic machine. The

toughness was measured at 4 K after each specimen was fatigue cracked in

liquid nitrogen at 76 K. Liquid nitrogen is less expensive than liquid

helium, and it ensures a much smaller plastic zone at the crack tip than the

alternative of precracking at room temperature. The maximum fatigue stress

intensity factor at the crack tip did not exceed 33 MPa-m^, the minimum- to

maximum stress intensity factor ratio, R, was constant at 0.1, and the loads

during fatigue were less than 40% of the limit loads for fracture at 4 K.

The value of Jj c was measured using the single -specimen unloading compliance

technique with computer-aided data reduction [4] , and Jj c was defined and

measured according to the ASTM E 813-81 method.

Fatigue crack growth rates were measured using the short crack simulation

described by Herman, Hertzberg, and Jaccard [5]

.

This differs from the

decreasing AK technique used in all previous cryogenic work [6-8]. In the

short crack simulation, the maximum fatigue stress intensity factor, Kmax , is

maintained constant while Km^n is progressively increased. The initial

conditions for this study are Kmax - 35 MPa-m^ and R - 0.1. AK is

subsequently reduced to near- threshold levels and R increases typically to

0.9. Using this technique for JK1 , we measured fatigue crack growth rates

spanning three orders of magnitude. Cycle frequency varied between 10 and 70

Hz depending on AK; higher frequencies were selected at lower AK to shorten

the time for measurements near the fatigue crack growth threshold. The

apparatus for fatigue crack growth threshold tests is described elsewhere

[ 7 ]

.

The procedure for calculating fatigue crack growth rates was as follows.

The crack was allowed to grow for an increment, typically 0.25 mm, and the

rate was calculated:

Aa/AN - (a2 -ai)/(N2 -N^)

.

( 1 )

The fatigue loads were then changed for measurements at a lower AK value

.

The new AK was calculated from [7]

AK - AK0 exp [C(a-a0 ) ]

,

( 2 )

where AK

AK0

a

aoC

instantaneous stress intensity factor range,

the initial stress intensity factor range,

the instantaneous crack length,

the final crack length, and

the stress intensity gradient,

rL " I 1

K da •

The selected value of C was typically 0.3 or 0.4 mm

-1

( 3 )

12
RESULTS

Tension Tests

The tensile property measurements at 295, 76, and 4 K are listed in

Table 2 (longitudinal orientations) and Table 3 (transverse orientations).

The data show that the JK1 conduit, before and after aging, has acceptable

ductility. The room temperature (295 K) ductility is slightly improved by

aging, whereas the cryogenic ductility (at 76 and 4 K) is slightly degraded.

The yield and ultimate tensile strengths increase strongly at cryogenic

temperatures as shown by the results for longitudinal specimens plotted in

Fig. 5. At 4 K, the yield strength exceeds 1200 MPa in both the as-received

and aged conditions. Aging systematically lowers the yield strength at each

temperature by a modest amount.

Table 2. Tensile property measurements for JK1 conduit specimens

in the longitudinal orientation.

Temperature

(K)

Condition Specimen

No

.

YS

(MPa)

UTS

(MPa)

EL

(%)

RA

(%)

E

(GPa)

295 As -received 1 . 889 987 30 65 188

2. 725 810 42 68 191

Avg. 807 898 36 66 190

Aged 1 . 680 932 29 75 174

2. 649 826 42 82 203

Avg. 664 879 36 78 188

76 As -received 1 . 1301 1555 51 64 201

2. 1351 1589 46 65 202

Avg. 1326 1572 48 64 202

Aged 1 . 1247 1555 33 43 204

2. 1096 1435 26 43 206

Avg. 1172 1445 30 43 205

4 As -received 1 . 1282 1713 49 48 200

2. 1427 1822 44 63 204

Avg. 1354 1768 46 56 202

Aged 1 . 1237 1780 37 53 200

2. 1738 38 53 198

Avg. 1237 1759 38 53 199

13
2000

JK1 Conduit

g. 1500

I

0

z

LD

DC

hCO 1000

YIELD STRENGTHS

ULTIMATE

STRENGTHS

500

1 00 200

TEMERATURE, K

300

Figure 5. Yield and ultimate strengths in the as-received (cold worked)

and aged conditions.

14
Table 3. Tensile property measurements for JK1 conduit specimens

in the transverse orientation.

Temperature

(K)

Condition Specimen

No.

YS

(MPa)

UTS

(MPa)

EL

(%)

RA

(%)

E

(GPa)

295 Aged 1 . NA 786 48 84 NA

2. NA 782 48 86 NA

Avg. 784 48 85

76 Aged 1 . NA 1426 62 84 NA

2. 1078 1442 65 85 217

Avg. 1434 64 84

4 Aged 1 . NA 1677 44 68 NA

2. NA 1651 47 67 NA

Avg. 1664 46 68

Fracture Toughness

Table 4 lists the Jjc measurements at 4 K. Tests in the T-L orientation

were performed in the as -received and aged conditions to show the effect of

heat treatment. Two tests were performed for each condition. The average

values at 4 K show that Jjc is reduced from 82 kJ/m2 in the as -received

condition to 40 kJ/m2 in the aged condition.

The effect of specimen orientation was evaluated (using aged material)

by comparing two tests in the T-L orientation and two in the L-T orientation.

Table 4 shows a measurable difference in the average Jjc values: 72 kJ/m2 for

the L-T orientation and 40 kJ/m2 for the T-L orientation. This is consistent

with the behavior of commercial rolled stainless steel plates in that L-T

orientations are usually tougher than T-L orientations.

The fracture surfaces of aged compact specimens tested at 4 K were

examined with a scanning electron microscope. In general, the fracture

surfaces are ductile and rough. For example, Fig. 6 features an aged

specimen in the T-L orientation tested at 4 K. The morphology consists of a

mixture of elongated troughs and small dimples. The troughs are oriented in

the same direction as the bands shown in Fig. 4. Figures 6A and 6B at

magnifications of 30 and lOOx show the zone of crack growth produced during

the J test. Elongated troughs extend from the fatigue crack front in the

direction of crack growth. Figures 6C and 6D show small dimples that formed

between the troughs. At 1000 and 3000x, rod-shaped inclusions are visible

near the center of the troughs. The same rod-shaped inclusions were cited in

Fig. 3. Energy Dispersive Analysis of X-rays (EDAX) indicates they may be

sulfides

.

13
Figure 6. Fracture surface appearance of JK1 compact specimen

in the T-L orientation after J- testing at 4 K:

(A) triangular zone of crack extension a (30x) , and

(B) enlargement, showing longitudinal troughs (lOOOx).

(C) enlargement, showing regions of fine dimples

adjacent to the troughs, along with rod- shaped

inclusions within the troughs (lOOOx).

(D) enlargement of above (3000x).

16
Table 4. Fracture toughness measurements for JK1 conduit at 4 K.

Specimen No. Orientation Condition J Ic

(kJ/m2

) Klc W*

(MPa -in

1

)

1 T-L As -received 120 155

2 T-L As -received 45 95

avg. 82 125

3 T-L Aged 38 88

4 T-L Aged 43 93

5 T-L Aged 40 90

avg. 40 90

6 L-T Aged 70 118

7 L-T Aged 75 122

avg. 72 120

a KIc (J) , an estimate of Kj c , is calculated from (E*Jj c/ •

Fatigue Crack Growth Tests

Fatigue crack growth rates for JK1 were measured at 295 and 4 K. Room

temperature results are shown in Fig. 7. The rates range from 10" 7 to 10" 10

m/cycle, for AK between 36 and 3 MPa-m^. There is no measurable effect of

specimen orientation (T-L vs. L-T) , judging from the data for material in the

aged condition. Based on data for the L-T orientation, aging does affect the

fatigue crack growth rates for this alloy, but only at relatively high AK.

At these higher values of AK (which correspond to low R ratios in this

study), the rates after aging are lower than those for as-received material.

At low AK (higher R ratios), the effect of heat treatment is either

nonexistent or too small to be detected beyond the scatter of measurements.

Therefore a single trend line at low AK in Fig. 7 represents all the data at

295 K. Following common convention, the threshold AKth can be defined

operationally as the value of AK that corresponds to a crack growth rate of

10" 10 m/cycle. Then, extrapolation of the data of Fig. 7 indicates that AKt^

is approximately 3 MPa*m^ for JK1 in both the as -received and aged conditions

at 295 K.

The fatigue crack growth rates measured at 4 K are plotted in Fig. 8.

As is the case at 295 K, there is no measurable effect of specimen

orientation (T-L vs. L-T) for specimens in the aged condition. Also, an

effect of heat treatment (as-received condition vs. aged condition) is

discemable only at AK higher than 10 MPa-m^. Below 10 MPa-m^, as threshold

is approached, the data for the as -received and aged conditions merge into

one trend line and extrapolation gives AKt^

<* 3 MPa'm^, which is equivalent

to the result obtained at 295 K.

17
Fatigue

Crack

Growth

Rate,

A3

/an,

m/cycle

Figure 7. Fatigue crack growth rates for JK1 conduit at 295 K.

18
Fatigue

Crack

Growth

Rate,

^/an,

m/cycle

Figure 8

.

Fatigue crack growth rates for JK1 conduit and other materials

at 4 K.

19
Figure 9 compares the data of the two previous figures, emphasizing the

effect of test temperature on the fatigue crack growth rates for this alloy.

The data are clearly temperature dependent at higher AK, but not at lower AK.

At higher AK where the temperature effect does exist, the rates at 4 K are

higher than at 295 K. At lower AK, below about 7 MPa-m^1

, the effect of

temperature becomes indistinguishable.

DISCUSSION

Static Mechanical Properties at 4 K

Figure 5 showed that heat treatment lowered the yield strength of the

as -received JK1 conduit. This is attributable to a recovery effect: heat

treatment relieved the plastic strain in the as -received conduit and

consistently reduced the yield strength at each temperature by about 140 MPa.

The aged alloy still meets the targeted strength of 1200 MPa or more at 4 K.

Figure 10 compares the conduit with plates tested previously [1,2].

Toughness is approximated by Kj c - (E-Jjc)^ and data for the T-L orientation

are compared. A Kj c value of 150 MPa -in* or more is desired for the conduit

[1], but the measured toughness is lower than that, and lower than levels

previously achieved by JK1 plates. Thus, the hoped-for properties were not

realized in the commercial product. Understanding the factors that determine

toughness is critical for product development, and factors influencing the

toughness are considered in the following discussion.

First, it is unlikely that specimen preparation or size can account for

the difference in toughness. The conduit specimens were full sized, and

machining tools never touched the planar surfaces. Our specimens were

thinner than the 25 mm plate specimens [1], but the Jjc measurements are

valid according to the Method E 813-81 size criterion. Moreover, the effect

of specimen size has been studied in stainless steels at 4 K. Typically,

thinner specimens give slightly higher, not lower toughness [9] so smaller

specimen size does not help to explain a lower measured toughness.

Differences in measurement techniques represent a potential influence on

Jj c results. However, a recent interlaboratory comparison revealed no bias

in the test techniques at the laboratories involved [10]. The older method

of data reduction (ASTM E 813-81) gives lower Jj c values than the current

method (ASTM E 813-88), simply because of a change in the definition of Jj c

. The Fig. 10 comparison, however, is based consistently on Method E 813-81.

The most significant factors appear to be material processing and heat

treatment. The data for plates in Fig. 10 represent tests performed on 90 kg

laboratory heats that were forged from ingots and hot- rolled at 1050°C into

28 -mm- thick plates. The plates were solution- treated and then aged at 700°C

for short (75 h) or long (200 h) time. Our specimens with a tube processing

history were aged at a lower temperature (650°C) for 180 h.

20
Figure 9 . Comparison of fatigue crack growth rate measurements for JK1

conduit using the short crack simulation technique (295 and 4 K)

.

The scatterband encompasses all measurements regardless of

ondition, specimen orientation, and temperature.

21
Figure 10. Strength—toughness comparison for JK1 conduit and plates in

relation to the trendline for commercial austenitic stainless

steel plates at 4 K.

22
Processing strongly influences the performance of JK1 . Even small

amounts of cold-working can be deleterious to properties after aging because

defects (dislocations) that accelerate diffusion and grain boundary

precipitation during subsequent aging are introduced [1]. Long aging times

are detrimental because they allows more time for precipitation. From plate

studies, the toughness after aging for 75 h is not much different than the

solution- treated condition; however, after aging for 200 h, the toughness

falls signifiantly and approaches the levels measured in this study for

conduit. Therefore, the lower toughness of JK1 in this study might be due to

a combination of residual cold work and a long (180 h) aging time.

From a metallurgical viewpoint, the mechanism of the toughness reduction

for the conduit remains unidentified. A more complete understanding would

require a separation of variables that is beyond the scope of the present

study. The complex JK1 microstructure features inclusions, precipitates,

and chemical inhomogenie ty. In principle, any or all of these factors can

contribute to a reduced toughness. We observed no grain boundary facets on

the 4-K fracture surfaces (Fig. 6), so apparently the level of grain boundary

precipitation after aging is insufficient to cause intergranular fracture in

this material at 4 K. The low toughness cannot, therefore, be ascribed

simply to grain boundary conditions. The rod-shaped inclusions might

contribute to low toughness if at relatively low strains they debond from the

matrix and produce voids.

Fatigue Crack Growth at 295 and 4 K

Figure 8 compared our measurements of fatigue crack growth rate with

Shimada's [2]. Shimada used the conventional decreasing AK procedure with a

constant stress ratio of 0.1. His results for JK1 plate are nearly

equivalent to NIST results for an AISI 316LN alloy [11] (see Fig. 8)

obtained by the same method. In this study of the conduit, we measured the

rates at constant Kmax and the stress ratio changed continuously from 0.1 to

0.9 as AK was lowered. Higher rates of fatigue crack growth are expected at

the higher stress ratios in this study. At AK - 36 MPa -in 1 and R - 0.1,

however, our rates for the conduit still exceed those for plates of JK1 and

AISI 316LN. We conclude that the rates of crack growth for the conduit in

this study are higher than those for AISI 316LN or JK1 plates because of

material as well as test differences.

Previously published fatigue crack growth rates for austenitic steels at

cryogenic temperatures apply for long cracks [6,7]. When the conventional

procedure is used for long- crack measurements near threshold the results at

low R are influenced by crack closure [6] whereas tests at high R are not.

Crack closure reduces the nominal AK applied, contributing to lower rates at

a given AK and higher thresholds at 10‘ 10 m/cycle. The technique used in the

present study ensures a high R ratio near the threshold. The AKtft measurements are then conservative, and the da/dN-vs.-AK data at 295 and 4 K

approximate short crack growth rates at these temperatures [5], It is

accepted that short cracks in structures will grow at rates faster than

predicted from long-crack data. Therefore, designers of superconducting

machinery may choose to base their design calculations on data produced by

the short crack simulation technique employed in this study.

23
SUMMARY AND CONCLUSIONS

A newly developed conduit sheath material was characterized with respect

to its mechanical properties. The material is 316LN, which contains 0.05%Nb

to retard carbide precitation during aging. The property measurements at 4 K

pertain directly to proposed ITER magnet designs.

The measurements of fatigue crack growth rate in this study used a

relatively new (constant Kmax) technique that can be used for conservative

design against failure by fatigue. We believe these data are the first of

their kind available for any material at 4 K; all previous cryogenic data

reported in the literature were obtained using the conventional decreasing AK

technique at a constant stress ratio, typically R - 0.1. Essentially, the

present results simulate the behavior of short cracks that are of practical

concern in conduit applications.

The yield strength of the JK1 conduit exceeds the targetted 1200 MPa at

4 K, but its fracture toughness (Jjc “ 40 kJ/m2

) is lower than anticipated

from earlier research and development reports. The results suggest that it

may be more difficult to achieve the hoped-for mechanical properties in

commercial products. Our experience here also underlines the importance of

establishing cryogenic material specifications and test standards, and

specifying fracture toughness minimums or verification testing for critical

materials in the procurement stage.

ACKNOWLEDGMENT

The test material was provided by N. Mitchell of the International

Thermonuclear Experimental Reactor and the NET research teams, Garching, West

Germany. The research was sponsored solely by the U.S. Department of Energy,

Office of Fusion Energy.

REFERENCES

1. M. Shimada and S. Tone, Effects of niobium on cryogenic mechanical

properties of aged stainless steels, in: Adv. Cryo. Eng., Vol. 34,

Plenum, New York, 1988, pp. 131-139.

2. M. Shimada, Fatigue crack growth rate at 4 K of aged austenitic

stainless steels, in: Adv. Cryo. Eng., Vol. 36, Plenum, New York, to be

published.

3. D.T. Read and R.L. Tobler, Mechanical property measurements at low

temperatures, in: Adv. Cryo. Eng., Vol. 28, Plenum, New York, 1982, pp.

17-28.

4. R.L. Tobler, D.T. Read, and R.P. Reed, Strength/toughness relationship

for strengthened AISI stainless steels at 4 K, in: Fracture Mechanics:

Thirteenth Conference, ASTM STP 743, R. Roberts, Ed., Amer. Soc. Test.

Maters., Philadelphia, 1981, pp. 250-268.

24
5 . W.A. Herman, R.W. Hertzberg, and R. Jaccard, A simplified laboratory

approach for the prediction of short crack behavior in engineering

structures, Fat. Fract. Eng. Maters. Structs., Vol. 11, No. 4, 1988,

pp. 303-320.

6. P.K. Liaw and W.A. Logsdon, Fatigue crack growth threshold at cryogenic

temperatures: a review, Eng. Fract. Mechs., Vol. 22, 1985, pp. 585-594.

7. R.L. Tobler and Y.W. Cheng, Automatic near- threshold fatigue crack

growth rate measurements at liquid helium temperature, Int. J. Fat.,

Vol. 7, No. 4, 1985, pp . 191-197.

8. R.L. Tobler and Y.W. Cheng, Midrange fatigue crack growth data

correlations for structural alloys at room and cryogenic temperatures,

in: Fatigue at Low Temperatures, ASTM STP 857, R.I. Stephens, Ed., Amer

Soc. Test. Maters., Philadelphia, 1985, pp. 3-30.

9. M. Shimada, R.L. Tobler, T. Shoji, and H. Takahashi, Size, sidegrooving, and fatigue precracking effects on J- integral test results fo

SUS 304 stainless steel at 4 K, in: Adv. Cryo. Eng., Vol. 34, Plenum,

New York, 1988, pp. 251-258.

10. H. Nakajima, K. Yoshida, S. Shimamoto, R.L. Tobler, P.T. Purtscher, and

R.P. Reed, Round robin tensile and fracture test results for an Fe22Mn-13Cr-5Ni austenitic stainless steel at 4 K, Adv. Cryo. Eng., Vol.

34, Plenum, New York, 1988, pp. 241-249.

11. R.L. Tobler, Near -Threshold Fatigue Crack Growth Rate Measurements for

Austenitic Stainless Steels at 4 K, in this volume.

25
'
NEAR -THRESHOLD FATIGUE CRACK GROWTH

OF AUSTENITIC STAINLESS

STEELS AT CRYOGENIC TEMPERATURES

R.L. Tobler

Materials Reliability Division

National Institute for Standards and Technology

Boulder, Colorado 80303

Measurements of fatigue crack growth rate are reported for 5

austenitic stainless steels at room and cryogenic temperatures (76

or 4 K) . The alloys evaluated in this study were AISI 316, 316LN,

304L, 304LN, and 304HN. The rate measurements span the range 10' 6

to 10" 10 m/cycle, and threshold stress intensity factors are

identified at 10" 10 m/cycle. The apparent thresholds for the 5

steels range between 6.5 and 10.5 MPa-m^ and vary slightly with

material and temperature. The effective thresholds (that is, the

apparent thresholds corrected for crack closure) are uniformly

lower than the apparent thresholds and range between 4 and 7

MPa-m^1

. At 4 K the thresholds increase moderately with yield

strength which ranges from 610 to 1270 MPa for these alloys.

INTRODUCTION

Fatigue crack growth rates and threshold stress intensity factors are

relevant to superconducting magnet design because they can be used to predict

safe operating conditions and lifetimes. Typically, fatigue crack growth

rates (da/dN) are plotted versus the stress intensity factor range (AK) ; the

data on log—log coordinates define a curve that may be divided into 3 stages:

very low rates approaching threshold values, intermediate rates of growth,

and high rates approaching instability.

In Stage I (low AK) the crack growth rates diminish and approach an

asymptote called the threshold, AKy^. Measurements near threshold are time

consuming and difficult to perform, so few data are available in this stage.

Liaw and Logsdon [1] reviewed these data in 1985. In Stage II (intermediate

AK) the measurements are easier and can be completed in relatively short

time; consequently the database for cryogenic materials in Stage II is larger

than in the other stages. Those data were reviewed by Tobler and Cheng [2],

also in 1985. In Stage III (high AK) , the crack growth rates for materials

accelerate as AK approaches the upper critical limit of fracture toughness.

This third stage is seldom studied because the crack growth rates are too

high for practical applications.

27
Owing to measurement difficulties the Stage I database for stainless

steels at cryogenic temperatures remains quite limited. All the problems

associated with room temperature testing, and more, are present in highcycle cryogenic testing. The cost and complexity of testing, as well as

material variability, is greater at cryogenic temperatures. Good

experimental technique is important, and automatic testing with computers to

acquire data in a reasonable time. Liaw and co-workers [3,4] designed a

fully automatic fatigue apparatus and report data for aged Fe-29Ni- 14Cr1 . 2Mo austenitic stainless steel and its welds at 297, 97, and 4 K. They

used a servohydraulic , computer-controlled test machine, conventional

fracture mechanics specimens, and a compliance technique. Later, we adopted

their method and evaluated 3 austenitic stainless steels at 4 K [5,6].

At the time of their review, Liaw and Logsdon found data for only one

austenitic stainless steel: AISI 304. Austenitic alloys, especially the

stainless steels, are important candidates for cryogenic magnet applications

in fusion energy. Therefore our purpose is to increase the data available

for these materials. This paper presents the initial results for long cracks

in five commercial 300 series stainless steels. Short crack behavior in

these alloys is the subject of a later paper.

PROCEDURES

Experimental Apparatus

The apparatus pictured in Fig. 1 is an automatic test system that allows

continous testing, data collection, and load adjustments to be performed by

computer without the assistance of a machine operator. Automatic testing

minimizes time and reduces data scatter, owing to higher precision in crack

length measurement and better control of data point spacing. The automatic

procedure is relatively easy to implement with minimal operator training, and

it eliminates subjective interpretation and operator bias.

The system consists of a 100-kN servohydraulic test machine and

cryostat, a programmable digital oscilloscope, a programmable arbitrary

waveform generator, a minicomputer and an automatic liquid helium refill

system. The servohydraulic machine has an inverted stage with its actuator

above the cryostat. The system includes a servo control system with two DC

conditioners, a valve drive, a signal amplifier, and a load cell. The

fatigue cycle counts are calculated by computer from the test frequency and

elapsed time. A programmable arbitrary -wave form generator is used since the

applied loads vary with time. The specimen is pinned at the bottom of the

load frame and enclosed during cryogenic tests in a vacuum- insulated

stainless steel Dewar. Nonmetallic materials of low thermal conductivity are

used to reduce the steady state heat leak into the dewar. The cryostat

design and associated instrumentation are described elsewhere in detail [6].

An automatic refill system, shown in Fig. 2, keeps the specimen fully

submerged in liquid helium for continous 24-h operation. The components

include a 305 -mm superconducting liquid- level detector, a level controller, a

conventional three-way solenoid value, a pressurized helium gas cylinder, and

28
a 500-1 liquid helium supply Dewar. A vacuum- jacketed flexible copper

transfer line connects the supply Dewar to the cryostat. The liquid helium

in the cryostat is monitored by the detector, which is attached to one of the

load frame columns. The desired minimum and maximum level limits are set on

the controller, and they trigger the solenoid valve that regulates liquid

helium transfer.

Figure 1. Computer-controlled fatigue test system.

29
Specimen Geometry and Clip Gage

The compact specimen used in this study is shown in Fig. 3. The

specimen width (W) is 76.2 mm and the thickness (B) is 6.4 mm. The

relatively large W/B ratio of 12 provides a sizeable ligament for crack

growth while minimizing the loads required. The machined notch is 22.9 mm

long (a/W - 0.3) and the notch tip radius is 0.05 mm. The clip gage used in

these tests is ring-shaped [7]. The gage was modified to use short beams

which enable the gage to be mounted tightly, close to the specimen. This

stabilizes the gage against vibration and reduces signal noise. Gage

calibrations at 295, 76, and 4 K are independent of temperature to within 2%.

Test Procedure

Crack length (a) is measured by compliance with the clip gage mounted at

the specimen edge. Compliance is specimen deflection per unit load, which

depends on crack length for a given material and specimen geometry. The

instantaneous crack length is computed using Saxena and Hudak's [8] function

to within 0.04 mm. The load cell and clip gage signals are interfaced to the

computer which infers the crack length once per minute and displays it on the

CRT terminal.

As the specimen is fatigued under the prescribed loading conditions and

cyclic frequency, a crack initiates and begins to grow. When the preselected

interval (1 min) for crack length measurement is reached, the computer

requests the compliance data from the oscilloscope, which instantaneously

freezes the load-vs . -displacement signals in memory and transmits them to the

computer. The computer fits the data to a straight line using a linear

least-squares method. Correlation coefficients of 0.995 or higher are

usually obtained.

Only a portion of the load-vs . -displacement curve (the part between a

specified intermediate load and a value corresponding to 95% of the maximum

fatigue load) is used for the compliance calculations. The specified

intermediate load is typically the mean load, because this is usually larger

than the crack closure loads. The upper 5% of the load is excluded to

eliminate any signal noise at maximum load during high-frequency tests.

As the crack grows, the difference between the current measured crack

length and the previously recorded crack length is checked against a

specified value (0.5 mm), which is within the recommended values of ASTM E

647-83. If the increment of growth is equal to or greater than the specified

value, the computer calculates a, N, da/dN, and AK. These data are stored on

floppy discs for later analyses.

In this study, the loads are progressively decreased according to the

decreasing-AK technique. The loads are calculated from the equations:

AK - AKp exp [ C { a-a0 } ]

,

( 1

)

AP - BW^AK/f (a/W) , and (2)

^max “ dP/(l-R), Pmin “ ^max^*

30
Figure 2. Schematic of automatic helium refill system.

Finish 2?. except holes, es noted.

(Dimensions in millimeters)

Figure 3. Fatigue test specimen.

31
where

Pmax - maximum load

Pmin " minimum load

R “ Rmin/^max

a - current crack length

a0 - crack length at the start of test

f (a/W) - [2 + (a/W)] [0.886 + 4.64(a/W) - 13.32(a/W) 2

+ 14 . 72 (a/W) 3 - 5.6(a/W) 4 ]/ [1 - (a/W)] 1 -* AK - current crack- tip stress intensity range

AK0 - crack- tip stress intensity range at the

start of test

C - a constant

The data were obtained in two steps. The rates from about 10' * to 10' 10

m/cycle were first measured by the decreasing AK technique starting at AK -

15 MPa-m1* with C - -0.08 mm' 1 and f - 40 Hz . After reaching AK-ph at 10 1 o The second data set was then measured by an

but with C

m/cycle, the test was stopped

increasing AK technique, starting again at about AK - 15 MPa-m

l 0.10 mm' 1 and f - 10 or 20 Hz. This second data set provided the rates up to

about 10' 6 m/cycle.

The sequence of operations is as follows. After each crack length

measurement, the crack length is compared with the last stored crack length

to ensure that a specified measurable amount of crack growth has occurred.

After the crack grows a certain amount (0.13 mm), the new AK is calculated.

At the relatively high frequencies desired in near- threshold fatigue

testing, hydraulic lag can occur. Then, the specimen is not actually

subjected to the load range commanded by the computer or waveform generator.

Overprogramming is sometimes necessary to overcome the problem. Automatic

overprogramming is a trial -and- error process whereby the computer monitors

the values of Pmax and Pm^n through the programmable digital oscilloscope and

makes necessary changes to achieve the commanded values.

During the test, point-to-point data reduction is used to calculate AK

and da/dN from a and N. If the results of da/dN-vs.-AK scatter excessively,

the seven-point incremental polynomial method can be used to smooth the

results after the test is completed.

For simplicity in this study the da/dN-vs.-AK curves and AK-j^ results

are presented on the basis of the applied AK values with no corrections for

crack closure. Since there are no accepted standard values of properties for

these materials, no meaningful statements can be made about the accuracy of

results. The measurement precision is commensurate with current practice.

Following the determination of AK-ph, a correction for crack closure is

applied. Crack closure causes the cracks to remain closed for a portion of

the load cycle; this portion of the load is ineffective in contributing to

the driving force for crack growth. In Fig. 4, the lower segments of the

bilinear load-displacement curves are associated with closure and are

disregarded in the calculation of the effective AK. The correction factor is

(Pmax ' Rc)/( pmax " pmin) » w^ere is the load at closure; this factor

multiplied times the apparent threshold gives the effective threshold.

32
Deflection,

0.008

mm/division

Deflection,

0.013

mm/division

Load, 0.2 kN/division

(A) T = 295 K,aK = 7.6 MPa m^

Load, 0.4 kN/division

(B) T = 4 K,aK = 8.5 MPa-m

1/2

Figure 4. Load-displacement records showing evidence of crack closure.

33
MATERIALS

Five austenitic stainless steels were tested. The materials were plates

and bars in the as-received, annealed conditions. AISI 316, 304L, and 304LN

were 51, 25, and 25 mm thick plates, respectively. Specimens from these

plates were machined in the T-L orientation according to the designation used

in ASTM Methods E 399-83. AISI 316LN was a 457 mm diameter bar, and AISI 304

HN was a 165 mm diameter bar; the specimen orientations for these alloys were

R-C and R-L, respectively. Tables 1 and 2 list the chemical compositions and

selected tensile properties for each material at room temperature and 4 K.

Table 1. Composition of test materials (mass %). Balance is Fe

.

Material Cr Ni C N Mn P S Si Cu Mo

AISI 316 17.25 13.48 0.057 0.030 1.86 0.024 0.019 0.58 • - 2.34

AISI 316LN 16.71 13.83 0.029 0.100 1.84 0.021 0.008 0.33 2.18

AISI 304L 18.20 8.56 0.020 0.082 1.73 0.024 0.005 0.62 0.24 0.40

AISI 304LN 18.00 9.91 0.037 0.120 1.50 0.022 0.017 0.57 0.197 0.205

AISI 304HN 20.21 8.44 0.059 0.28 1.83 0.025 0.021 0.50 0.60 0.59

Table 2. Measured mechanical properties of test materials.

Material Test Yield Ultimate Elongation Reduction

temperature strength strength of area

(K) (MPa) (MPa) (%) (%)

AISI 316 295 228 576 56 73

4 711 1301 48 57

AISI 316LN 295a 267 551 53 76

4 834 1364 45 53

AISI 304L 295 245 642 66 77

4 610 1678 35 48

AISI 304LN 295 279 612 68 62

4 745 1570 34 42

AISI 304HN 295 a 441 765 47 73

4 1267 1727 8 22

aRoom temperature data are from manufacturer's mill sheet.

34
RESULTS

The da/dN-vs.-AK curves for the 5 stainless steels at selected

temperatures and stress ratios are presented in Figs. 5-13. In these

figures, da/dN is plotted versus the applied AK values without closure

corrections. Following conventional methodology, the threshold is

operationally defined as the value of AK corresponding to an average rate of

10" 10 m/cycle. This rate corresponds to less than 1 lattice spacing per load

cycle. The apparent thresholds (to the nearset 0.5 MPa.mH) were obtained

from the data plots and are listed in Tables 3 and 4, along with the

effective thresholds which are from the apparent thresholds with corrections

for crack closure.

AISI 316

Figure 5 shows room temperature results for AISI 316 at R - 0.1. The

data trend on log—log coordinates is linear, and the rates at low AK do not

begin to approach an asymptote at 10" 10 m/cycle. The measurements agree well

with rates reported by Read and Reed [9], who tested this same heat in Stage

II at AK - 25 to 50 MPa-m^. Our data at low AK complement their data for 25-

mm thick compact specimens at 20 Hz. This indicates there is no effect of

specimen thickness between 6 and 25 mm, that the Stage II linear trend for

this alloy apparently extends to exceptionally low growth rates, and that

true thresholds, if they exist, could be detected only by measurements at

lower AK values than studied here.

The 4-K measurements for AISI 316 at R - 0.1 are also shown in Fig. 5.

Again, the rates are in fair agreement with the previous Stage II results

[10] for 25-mm thick specimens, but the data trend differs from that at 295

K. Over the broader range of AK the da/dN curve at 4 K shows a knee and

crosses the 295-K trend line twice. Clearly, for AISI 316 at R - 0.1,

reducing the temperature from 295 to 4 K increases the threshold for fatigue

crack growth. The apparent AK-ph at 4 K (8.5 MPa-m

1*) is slightly higher than

that at 295 K (7.5 MPa-m^1 ). When crack closure corrections are applied, the

effect of temperature on AK^h is maintained, as shown by the effective

threshold values listed in Table 3.

Fatigue crack growth at a higher stress ratio, R - 0.3, is shown in Fig.

6. This time the apparent AK*^ decreases as the test temperature is

decreased from 295 to 4 K. This is the opposite of the temperature

dependence shown in Fig. 5 and Table 3 for apparent thresholds at R - 0.1.

As explained later this apparent reversal of the temperature dependence is

due to crack closure. The effective values of AK^h do follow the same

temperature dependence at both R ratios, as shown in Table 3.

By comparing the room temperature data in Figs . 5 and 6 , we see that the

295 K rates and thresholds are independent of stress ratio at R - 0.1 and

0.3. At 4 K, however, a significant effect of R ratio emerges, as shown by

the data in Fig. 7. In tests at 4 K, we examined the behavior at 4 different

stress ratios between 0.1 and 0.7. For the range of AK studied, the 4-K

rates are equivalent at R - 0.3, 0.5, and 0.7, but higher than at R - 0.1.

In addition, the range of Stage II is dpressed at the higher R ratios as

compared to R - 0.1. AK-^ is also lower at R - 0.3, 0.5, and 0.7 than it is

at 0.1.

35
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle

Stress Intensity Factor Range, AK, MPa-m/2

Figure 5. Fatigue crack growth rate data for AISI 316 at R - 0.1.

36
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle

1 10 100

Stress Intensity Factor Range, AK, MPa*m/2

Figure 6. Fatigue crack growth rate data for AISI 316 at R - 0.3.

37
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle

Stress Intensity Factor Range, AK, MPa- m

'2

Figure 7. Fatigue crack growth rate data for AISI 316 at several

stress ratios.

38
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle

Figure 8. Fatigue crack growth rate data for AISI 316LN at R - 0.1

39
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle

Stress Intensity Factor Range, AK, MPa*m/2

Figure 9. Fatigue crack growth rate data for AISI 304L at R - 0.1.

40
Figure 10. Fatigue crack growth rate data for AISI 304LN at R - 0.1.

41
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle

Figure 11. Fatigue crack growth rate data for AISI 304HN at R - 0.1.

42
10

Austenitic Stainless Steels

Figure 12. Fatigue crack growth rate data for 5 stainless steels at 4 K.

43
Threshold

Stress

Intensity

Factor

Range,

AK

Th

,

MP-m^

Threshold

Stress

Intensity

Factor

Range,

AKp,,

MP-m

Figure 13. Dependence of fatigue crack growth thresholds on

yield strength.

44
Table 3. Apparent and effective threshold stress intensity factors for AISI

316 stainless steel at two temperatures and two stress ratios.

Temperature

(K)

Stress Ratio, R Apparent

(MPa-m1*)

Effective Krh

(MPa-m1*)

295 0.1 7.5 4.6

0.3 7.5 4.6

4 0.1 8.5 6.1

0.3 6.5 5.1

Table 4. Apparent and effective threshold stress intensity factors

for 5 austentitic stainless steels at R - 0.1.

Alloy Temperature Apparent Threshold, K^h Effective Threshold

AISI 316 295 7.5 4 .

6

4 8.5 6.1

AISI 316LN 4 8.5 6.8

AISI 304L 295 6.5 4.3

4 7.5 4.4

AISI 304LN 295 9.0 4.0

76 8.0 4.4

4 8.0 6.3

AISI 304HN 295 9.0 4.0

76 10.0 7.0

4 10.5 6.8

45
AISI 316LN

AISI 316LN contains less carbon and more nitrogen than AISI 316. The

yield strength of AISI 316LN is consequently higher at 4 K but otherwise the

two alloys have nearly equivalent mechanical properties in tension tests

(Table 2). Fatigue crack growth rates for AISI 316LN were measured in liquid

helium at R - 0.1, and the results are plotted in Fig. 8. The data define a

curve similar to that for 316. The apparent threshold, 8.5 MPa-m^, is equal

to that for AISI 316 at T 4 K and R — 0.1. The AISI 316LN alloy was not

studied further since it behaves similarly to AISI 316.

AISI 304L

Figure 9 presents the data for AISI 304L. The 304L (and other 304- type

steels) has substantially less Ni than 316 or 316LN; it also has a lower

yield strength but a higher ultimate strength than 316 or 316LN. The fatigue

crack growth rates for 304L at R - 0.1 were measured at 295 and 4 K and the

results are shown in Fig. 9. In general, the 304L behavior is similar to

that of AISI 316 and 316LN at the conditions studied: the 295 K data trend is

nearly linear, the 4 K trend is sigmoidal, and the two trend lines cross

twice. The AK*^ are 6.5 MPa-rn^1 at 295 K and 7.5 MPa-m^ at 4 K.

AISI 304LN

The 304LN steel of this study has 0.04% more N than the 304L, which

makes it stronger than the 304L, and it has 2% more Ni which makes it more

stable against phase transformation. Figure 10 presents the data for AISI

304LN at R - 0.1. The fatigue crack growth rates were measured at 295, 76,

and 4 K. The rates for 304LN are similar to those for 304L over most of the

AK range studied. The data at 295 K do not show an obvious knee at low AK.

A gentle curve suffices to describe the trend. In stage I, the 76 and 4 K

results appear to be essentially equivalent. The apparent thresholds at both

these cryogenic temperatures are 8 MPa-m^ which is lower than the 9.0 MPa-m^5 value obtained at 295 K. However, the effective threshold is higher at 4 K

than at room temperature or 76 K.

AISI 304HN

AISI 304HN contains more N than any other alloy of this study. As a

result, the alloy has greater strength and lower ductility than the others.

The yield strength of the 304HN at 4 K is particularly high and exceeds 1200

MPa. Figure 11 presents the measurements of fatigue crack growth rates for

AISI 304HN at three temperatures (295, 76, and 4 K) and R - 0.1. The

apparent thresholds at each temperature are the highest reported for the 5

steels tested in this study. Again the apparent and effective thresholds

increase at cryogenic temperatures. The results at 76 and 4 K are

essentially equivalent, and higher than at 295 K. At 4 K the apparent and

effective thresholds reach 10.5 and 6.8 MPa-m*1

, respectively, and are clearly

the highest values observed for any alloy in this study.

46
Comparison

Figure 12 compares all 5 alloys at T - 4 K and R - 0.1. Overall, the

data define similarly shaped curves that fit in a relatively narrow band,

with appparent thresholds ranging from 6.5 to 10.5 MPa-m^. Some

metallurgical effects are perceived in the results. For example, in stage It

at AK - 15 MPa-m^5 or more, the rates become higher for the alloys in the 304

system than in the 316 system. Also, at the highest values of AK studied,

the strength of alloys may be influential within each system since the rates

for 316LN are lower than for 316, and the rates for 304HN are lower than for

304L or 304LN. As AK decreases to intermediate values, the curves for

different alloys cross over. As threshold is approached, the yield strength

influences the ranking of alloys. This is shown clearly in Fig. 13, where

the thresholds (both apparent and effective) are plotted versus yield

strength. At 295 K, the trend is insignificant because there is not a great

variation in yield strengths for the 5 steels at this temperature. The

effective thresholds at this temperature are equivalent for the 5 alloys. At

4 K, the dependence on yield strength is more pronounced. As shown in this

figure, the apparent and effective thresholds both increase moderately with

yield strength at 4 K.

DISCUSSION

Liaw and Logsdon's [1] review considered cryogenic threshold data

available at the time for aluminum, copper, nickel, steels, and titanium

alloys. (Only one 300 series steel, AISI 304, was referenced.) They

observed that decreasing temperature generally increases the threshold. They

used crack closure concepts to explain the effects of stress ratio and

dislocation dynamics to explain the improvement of thresholds at cryogenic

temperatures. The same principles are sufficient to rationalize the behavior

of the 300 series stainless steels in the present study.

At 295 K, as shown in Fig. 5, the fatigue crack growth rates for AISI

316 follow a linear trend for at least four orders of magnitude in da/dN, and

no asymptote is approached at 10“ 10 m/cycle. The 304L and 304LN alloys

behaved similarly at this temperature, as shown in Figs. 9 and 10. This

observation implies that no true threshold is exhibited in this range, and

the operational definition of AK^ at 10" 10 m/cycle is evidently of limited

significance for the austenitic stainless steels at room temperature.

Presumably, crack growth could be measured at AK below the operationally

defined thresholds, although the time required for such measurements by

present methods would greatly increase. An average growth rate of 10" 10

m/cycle corresponds to 1 angstrom per cycle whereas the lattice spacings for

austenitic steels are a few Angstroms. Rates of crack growth less than a

lattice parameter per cycle are possible, assuming that the crack grows

nonuniformly , in spurts, or in highly localized areas along the crack front.

Additional tests at still lower AK values are needed to confirm that crack

growth continues at rates below 10" 10 m/cycle for other austenitic alloys.

47
Stress Ratio and Crack Closure Effects

Crack closure as observed in the oscilloscope load—displacement signals

figures prominently in the interpretation of our results. Closure is more

pronounced in the austenitic steels at room temperature than at cryogenic

temperatures. It is also more pronounced at low AK than at high AK. In the

tests at 295 K near threshold, closure loads typically approached the mean

load of the fatigue cycle at R - 0.1, whereas at 4 K the closure loads were

much lower due to increased resistance to plastic deformation.

For AISI 316 between R - 0.1 and 0.3 we observed an effect of stress

ratio at 4 K but not at 295 K. The closure load was significantly higher at

295 K than at 4 K, as shown in Fig. 3. In this case the closure and minimum

loads were such that the change in stress ratio from 0.1 to 0.3 affected the

effective AK at 4 K, but not at 295 K.

Crack closure also explains why there is no effect of stress ratio

between R - 0.1 and 0.3 for AISI 316 at 295 K. At 295 K the closure load is

about h the applied maximum loads; therefore the effective AK is the same at

R - 0.1 and R - 0.3 (Table 3), and no stress ratio effect is observed.

Stress ratio effects would be expected at R ratios above 0.5.

Crack closure also explains the effects of stress ratio found at low R

at 4 K. At 4 K, rate differences are measureable between R - 0.1 and 0.3

(Fig. 6) because closure occurs at low loads; the closure load is

intermediate between the minimum loads at R - 0.1 and 0.3 and that affects

the effective AK.

Temperature Effects

Owing to closure the apparent thresholds for these alloys do not always

increase with a temperature reduction from 295 to 4 K. However, when closure

is taken into account, the effective thresholds are always higher at 4 K than

at 295 K. Other metals and structural alloys at room and cryogenic

temperatures show the same result: the effective threshold at cryogenic

temperature is higher than at room temperature which is consistent with

dislocation dynamics theory [1].

In general, the steels in this study behaved similarly at 4 K (Fig. 9).

One reason for this is that all have similar Young's moduli. We showed that

rates for a wide variety of alloys in Stage II depend strongly on Young's

modulus and can be normalized on the basis of the parameter AK/E [2]. In

Stage I also, the threshold increases with Young's modulus [10]. Young's

modulus is fixed primarily by the base composition and is structure

insensitive. The 5 alloys in this study are composed primarily of iron,

chromium, and nickel, and their Young's moduli are equivalent within a few

percent. The general similarity of performance is therefore not surprising.

Aside from Young's modulus, other factors influence fatigue resistance.

The secondary effects of yield strength were cited in Figs. 12 and 13.

Austenite phase stability is potentially important. In very unstable alloys,

48
especially those with low C, N, and Ni contents, massive martensitic phase

transformations can improve fatigue crack growth properties [11,12]. The

AISI 304L steel of this study has low C but fairly high Ni and other

stabilizing elements. Additional tests of still leaner 304L grade alloys are

recommended to evaluate possible effects on fatigue crack growth. However,

even if the effect proves to be favorable, 304L grades would not be used for

fusion magnets at 4 K since high yield strength is required. High yield

strengths in austenitic stainless steels is achieved primarily by nitrogen

additions which also stabilize the austenite against phase transformation.

The data in this study pertain to long cracks in conventional compact

laboratory specimens. To date, most cryogenic data for stages I and II have

been produced for specimens containing long cracks, typically at low R, and

frequencies from 4—20 Hz. The use of these data may not be appropriate in

applications where short cracks are present. Room temperature studies show

that short cracks do not behave like long cracks. Short cracks grow at

higher rates than data from specimens with long cracks might suggest. To

support proposed superconducting machinery designs, both long and short crack

behavior must be characterized and understood. Therefore short crack

behavior is being simulated in ongoing research at NIST and is the subject of

a subsequent paper.

SUMMARY AND CONCLUSIONS

The fatigue apparatus used in this study is fully automatic. The

approach allows continuous data reduction and load adjustments without an

operator. Tests at 40 Hz last 4 days and nights with continuous 24-h

operation. Typically, 13 million load cycles are needed for 1 threshold

measurement and to measure rates of crack growth spanning the range between

10" 6 and 10" 10 m/cycle. Fatigue crack growth rates for 5 austenitic

stainless steels are reported. The decreasing-AK technique was used at

selected R ratios to obtain data for long cracks in laboratory specimens.

Several material and test variable effects are summarized as follows:

1. The rates at 295 K for several austenitic stainless steels do not

approach an asymptote at 10" 10 m/cycle; thus crack growth will occur at lower

rates, and longer tests may be needed to reach a true threshold at 295 K.

2. The effect of a temperature reduction from 295 to 4 K is to increase

the effective thresholds of all 5 steels.

3. At 4 K, conventional thresholds exist near 10" 10 m/cycle for steels

tested. The operationally defined thresholds increase slightly as yield

strength increases.

4. At 4 K, the apparent thresholds range between 6.5 and 10.5 MPa*m . The effective thresholds (the apparent thresholds corrected for crack

closure) are uniformly lower and range between 4 and 7 MPa-m^.

49
REFERENCES

1. P.K. Liaw, and W.A. Logsdon, Fatigue crack growth threshold at cryogenic

temperatures: a review, Eng. Fract. Mechs., Vol. 22, 1985, pp. 585-594.

2. R.L. Tobler and Y.W. Cheng, Midrange fatigue crack growth data

correlations for structural alloys at room and cryogenic temperatures,

in: Fatigue at Low Temperatures, ASTM STP 857, R.I. Stephens, Ed., Amer.

Soc. Test. Maters., Philadelphia, 1985, pp. 3-30.

3. P.K. Liaw, W.A. Logsdon, and M.H. Attaar, Automated near- threshold

fatigue crack growth rate testing of JBK-75 stainless steel at cryogenic

temperatures, in: Austenitic Steels at Low Temperatures, R.P. Reed and

T. Horichi, Eds., Plenum, New York, 1983, p. 171-185.

4. P.K. Liaw, W.A. Logsdon, and M.H. Attaar, Cryogenic temperature nearthreshold fatigue crack growth rate data for JBK-75 stainless steel,

Cryogenics, Vol., 23, 1983, pp. 523-526.

5. R.L. Tobler, and Y.W. Cheng, Automatic near- threshold fatigue crack

growth rate measurements at liquid helium temperature, Int. J. Fat.,

Vol. 7, 1985, pp. 191-197.

6. R.L. Tobler, Near- threshold fatigue crack growth behavior of AISI 316

Stainless Steel, Adv. Cryo. Eng., Vol. 32, Plenum, New York, 1986, pp.

321-327.

7. R.L. Tobler and J.A. Shepic, Design and performance of a ring-shaped

clip gage for fracture mechanics testing, J. Test, and Eval., Vol. 13,

No. 4, 1985, pp. 299-302.

8. A. Saxena and S.J. Hudak, Jr., Review and extension of compliance

information for common crack growth specimens, Int. J. Fracture, Vol.

14, No. 5, 1978, pp. 453-468.

9. D.T. Read and R.P. Reed, Fracture and strength properties of selected

austenitic stainless steels at cryogenic temperatures, in Materials

Studies For Magnetic Fusion Energy Applications at Low Temperatures -

II, NBSIR 79-1609, National Institute of Standards and Technology,

Boulder Co, 1979, pp. 81-122.

10. P.K. Liaw, W.A. Logsdon, and M.A. Burke, Fatigue crack propagation

behavior of 63Sn-37Pb solder. Scripts Metall. , Vol. 23, 1989, p. 747.

11. Y. Katz, A. Bussiba, and H. Mathias, The role of induced phase

transformations on fatigue processes in AISI 304L, Proceedings of the

4th European Conference on Fracture, p. 503.

12. Mei, Z. , Chang, G.M., and Morris, J.W., Jr., Fatigue crack growth in

metastable austenitic stainless steels, in: Cryogenic Materials, 88,

Vol. 2, Structural Materials, R.P. Reed, X.X. Xing, and E.W. Collings,

Eds., International Cryogenic Maters. Conference, Boulder, Co., 1988,

pp. 491-500.

50
CREEP OF COPPER: 4 to 295 K

R.P. Reed, N.J. Simon, and R.P. Walsh

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado 80303

ABSTRACT

Creep measurements at 295, 76 and 4 K have been conducted on C10400

copper. Specimens were held under constant tensile load (dead weight) for

periods of time sometimes exceeding one month. Creep data have been fitted

to a series of expressions relating creep strain to time by a nonlinear

least-squares procedure. The coefficients of the terms in these equations

were related to applied stress levels. Primary, logarithmic, and steadystate creep ranges are discussed.

INTRODUCTION

For many cryogenic technological applications, such as the operation of

a high- field superconducting magnet for long periods of time, knowledge of

the creep properties of the structural reinforcements is necessary. The

study of the time -dependent mechanical behavior of structural materials at

low temperatures has proved very challenging. Maintenance of temperature

and mechanical stability, as well as accurate recording of stress and strain

over long periods of time, stretch the limits of our measurement capabilities .

The creep of metals at low temperature has been reviewed by Tien and

Yen 1 and Yen et al. 2 A wide range of results are apparent from these

reviews. Both logarithmic (creep strain proportional to logarithm of time)

and steady-state (creep strain proportional to time) have been reported for

metals at low temperatures.

Copper is used extensively as a stabilizing material for superconducting magnets and as a conductor in high field, normal -metal magnets.

Thus, it is necessary to understand its creep characteristics. The creep

of copper, perhaps, has been studied more chan any ocher metal at low

temperatures. 1 * 7 This paper reports on the creep of copper at temperatures

less chan 0.22 Tm (295, 76, and 4 K) and at stress levels about 10* 3 ^ (25 -

60 MPa) where Tra is the melting temperature (1356 K) and m is the shear

modulus. Our study provides a basis for comparison of the results from

various laboratories and for understanding the complexities of these

measurements

.

Advances in Cryogenic Engineering (Materials/. Vol 36 1175

Edited bv R P Reed and F R. Pickett Plenum Press. Ne* York. 1990

51
Table 1. Creep Strain Expressions

(1) <erue “

<o * a iln t (7)

(2) “ <o + a i^ t + a5 t (8)

(3) <true “

<o + a iln(a,t +• 1) (9)

(4) «enje “ <o ajlnCajC + 1) a5 t (10)

(5) <tzue <o + ai< 1 * as c (11)

(6) ‘true “ <o + a iC l/3 + a,t2 /3 + as t

‘true

‘true

‘true

‘true

‘true

<o + + a*t

< 0 + a^* a5 r4* + a5 t

<o + a,ln(a,c a3 )

<o + ajln^t + as ) a^c

< 0ajln c + ajt**

EXPERIMENTAL PROCEDURES

Oxygen-free, high conductivity C10400 copper (Cu + Ag - 99.99 vt.%) wa obtained in 20-mra diameter bar stock. It was used in the annealed conditio

(650 # C/1 h) with a hardness of Rg - 22 and a grain size of 36 pm. Subsequently, specimens were reannealed (650*C/1 h) after undergoing creep

deformation of 0.02 to 0.04 in previous tests; these specimens are called

"reannealed" in the text. Average yield strengths of 25. 27, and 30 MPa

were measured at 295, 76 and 4 K, respectively, for this copper.

The dead-weight loading system utilized a lever arm pivoted above the

specimen. For low temperature tests a 900-mra deep superinsulated dewar with

a narrow neck was used. Dewar capacity was 30 L. The test fixture con- sisted of a G-10CR outer compression cylinder and an inner titanium pull

rod. Liquid helium boiloff without the test fixture was 0.2 L/h and with

the test fixture was 0.5 L/h.

Metal-film, resistance strain gages (73Ni-20Cr alloy) were bonded to

the round specimens with a low- temperature epoxy. Strain sensitivity at room and low temperature was 5 x 1 0

7

; instrumental and temperature drift

resulted in strain readout variability over a period of one month of about

± 2 x 10* 7 at room temperature and about ± 4 x 10* # at 4 K. To measure variability during a test, dummy specimens were inserted in the cryostat.

Other environmental factors caused measurement inconsistencies . These

included vibrations, ambient temperature variability, cryogenic liquid

transfer, and the reduction of the cryogenic liquid level with time. Floor

vibrations induced by ocher laboratory equipment led to sudden spurious

jumps (- 10* s

) of strain during testing. Vibrations induced specimen strain

despite the use of damping pads. Ambient temperature variability led to

daily cyclic patterns of strain (at the 10* 7 level); to eliminate these

thermally induced strain cycles subsequent test specimens were maintained at

303 K using resistance heaters attached to brass specimen grips. The

transfer of cryogenic fluids into the dewar produced transient specimen

loads of the order of 0.5 kg chat were induced by the differencial thermal

contraction between the pull rod and compression cylinder. These transient

loads, measured by strain gages on the pull-rod and compression cylinder, were not adequately compensated by the lever-arm. Similarly, the reduction

of liquid level with time is thought to have led to load changes in the

filament-reinforced composite cylinder from thermal contraction changes

transmitted by the continuous filaments. These strain differences, again in

the range of 10* 7

, led to small uncompensared load changes on the specimen.

MODEL ANALYSIS

The creep strain after a period of elapsed time, t, is the sum of the

initial strain upon loading, < 0 , and a time -dependent strain. Both the

1 1 76

52
Table 2. Summary of Creep Daca

Ttap

.

(K)

° a <«?•)

°*/°y *o (10**)

»» (10*‘)

Durtelon

(Bln)

295 20.20 0.81 1.704 3.192 37440

24.90 1.00 3.269 4.153 37440

27.00 1.08 0.709 1.67 29986

29.96 1.20 4.726 3.44 27300

31.40 1.26 4.873 3.15 18856

34.00 1.36 3.131 3.663 29986

34.10 1.36 5.575 4.254 41640

36.40 1.46 5.647 4.297 49950

37.60 1.50 7.031 4.871 41640

31.00 1.52 4.083 4.371 29986

41.00 1.64 5.596 5.111 29986

42.20 1.69 8.208 5.845 13695

42 . 40 1 . 70 8.031 4.232 49950

45.10 1.80 9.116 4.49 18856

45.30 1.81 9.032 5.672 13695

76 24.30 0.90 0.650 2.323 12591

26.10 0.97 1.659 3.233 12591

29.70 1.10 2.609 2.306 96600

33.20 1.23 2.23 3.934 96600

35.20 1.30 1.181 3.763 12591

35.40 1.31 2.731 3.876 96600

37.10 1.37 4.889 4.684 58894

41.30 1.53 4.815 3.863 36345

45 00 1.67 5. 251 4.388 43435

49.50 1.83 7.499 4.642 12820

u 30.00 1.00 2.199 0.003 1027

30.60 1.02 0.838 0.361 21339

32.00 1.07 3.521 0.139 1027

34.40 1.15 2.563 0.663 21339

35.60 1.19 3.817 1.008 1488

37.60 1.25 2.171 0.423 21339

41.50 1.38 5.743 0.787 18540

43.80 1.46 4.933 0.16 3612

44.00 1.47 5.834 0.17 1027

47.20 1.57 3.926 0.144 1403

48.90 1.63 8.223 1.166 6935

60.00 2.00 17.56 2.283 6935

initial and cime -dependent scrain components are functions of the temperature, T, and the applied stress, o a

. We have chosen to describe this

functional dependence by first fitting a series of expressions of the creep

scrain as a function of time (Table 1) to individual daca secs obtained at

4, 76, and 295 K for a series of <7 a values. Then, the dependence of the

constants (a^) in these expressions upon o a and T was determined from plots

of the constants versus o a for the three test temperatures. The 11 expressions in Table 1 were obtained from the literature, or are generalizations

of expressions for creep scrain found in the literature. The creep scrain

expression is expected to be logarithmic [Equations (1) and (3)] when

T 5 0.3 Tra Equations (2) and (4) allow for a steady-state creep (a 5 c,

e - a 5 ) in addition to the logarithmic creep, and Eqs. (9), (10), and

(11) are further generalizations of these expressions. Equations (5) and

(6) are expressions given in the literature for temperatures of the order of

0.4 Tm and higher. In this temperature regime, the primary stage of creep

is represented by [a t (1 - e* a 2 ^)] or [ait 1 / 3 + a 2 t

2 ' 3

] and the secondary,

or steady-state regime, by the a s c term. Equations (7) and (8) are generalizations of these expressions. The equations in Table 1 were fitted to the

daca secs by a nonlinear least- squares computer program. The constants a^

from the expressions chat best fitted the daca secs were then plotted as a function of a a

. A linear least - squares fit was made to the daca in these

plots

.

1 1 77

53
EXPERIMENTAL RESULTS AND DISCUSSION

The results are summarized in Table 2, which incLudes applied stress,

the ratio of the applied stress to the yield strength (<7y) , the initial

strain on application of the dead weight (< 0 ), and the coefficient for the

primary creep term (a t ) from Eq. (2) in Table 1. When the data sets at 295

and 76 K were fitted to the equations in Table 1, Eq. (11) usually gave the

lowest standard deviation (S.D.). In many cases, the nonlinear least

squares procedure did not give a convergent result when the data set was fitted to Eq. (11); however, the S.D. of the fit was the lowest found for 12

out of 15 sets at 295 K, and for 6 out of 7 sets at 76 K (Eq. (11) was not

fitted to 3 of the 76-K data sets reported in Table 2). The test duration

ranged from 15 700 to 50 000 min at 295 K and from 12 600 to 96 600 min at

76 K. At 4 K, the test duration was generally much shorter (1000 to 21 300

min) and Eq . (11) was not found to give the best fit to the data sets. It

is thought that Eq . (11) usually provided the best fit for data secs of

longer duration because the power of t(a 4 ) was allowed to vary; in Eqs. (2),

(4) - (8), and (10), the power of t is 1.0. At the longer elapsed times, it

is more likely that a different stage of creep has been entered, and

therefore, a more general expression for t allows a better fit to the data.

[That multiple stages of creep were present is further substantiated by the

observation that in all cases. Eq . (2) was a better fit than was Eq. (1) and

Eq. (4) was better than Eq . ( 3 )

]

.

However, because small aberrations in a data sec could exert a large influence over the values of a, and a 5 for Eq.

(11) determined in the nonlinear least-squares iterations (a t is slightly

affected, also) the functional dependence of the constants e Q and a t upon T

and a a for Eq . (2) are used in our discussions.

Instantaneous Strain

On application of load there is an instantaneous specimen strain. The

square root of the initial strain, « 0 , is a linear function of the applied

stress, a a

. This is illustrated for all temperatures in Fig. 1 and sum- marized in Table 2. This result is expected, since the stress-strain curves

of copper are approximately parabolic and the primary strain on application

of load corresponds to this relationship.

Fig. 1. Applied stress versus specimen strain (to 1/2 power) for copper at

295, 76. and 4 K.

1 1 78

54
8600

8500

8400 f*

o

Z 8300

2

CO 8200

8100|-

8000

4 K

«•, s 48 9 MPa

10 100

TIME. MIN

1000

Fig. 2. Dependence of creep strain versus log time for copper under a stress

of 48.9 MPa at 4 K.

Creep Behavior: General

Tests were conducted at three temperatures (4, 76, and 295 K) and at

relatively low stress levels (- 10* 3 m). When the test temperature is less

than approximately 0.3 Tra , logarithmic or exhaustive creep is normally

expected. (Here we define exhaustive creep as distinguished by an ever- decreasing creep rate, where the creep strain is not a linear function of

the log of the time.) In our creep tests the test temperature is less chan

0.22 Tm. Steady-state creep races or a normal secondary creep stage are not expected to occur under these test conditions.* Yet, there have been a number of papers attempting to analyze for activation energies and to measure steady-state creep races, even at 77 K. These include the recent

studies by Yen et al. 3,5 and reviews by Tien and Yen and Yen ec al. 2 emphasizing thermally activated steady-state creep processes in copper at

77 K. We found distinct creep characteristics of copper at each test

temperature. However, in no case was steady-state creep detected. Results

are presented and discussed separately for each test temperature.

fig6

5

4

r2 3 n

2

1

20 30 40 50 60 70 STRESS l<r,). MPa

, versus applied stress from Eq.

76 and 4 K.

0 10

3. Creep coefficient,

for copper at 295,

2 (Table 1)

1 1 79
5400

5300

5200

5100

5000

4900

76 K

». = 41 3 MPa

4800 h

10 100 1000

TIME, MIN

10000

Fig. 4. Dependence of creep strain on Log time for copper under a stress of

41 . 3 MPa at 76 K.

Creep Behavior: 4 K

Initially, the creep strain at 4 K is linearly related to the log of

time. But after a brief time (usually less than 1000 min) the dependence of

strain on log time decreases. This dependence is illustrated in Fig. 2.

The logarithmic coefficient, a t

, from Eq . (2) (Table 1) is plotted versus applied stress ( a a ) for 4 K in Fig. 3. Compared to higher temperature

coefficients. is quite low. The data of Yen et al. 3 at 4 K correspond

very closely to our data where a x is of the order of 10‘ s. Startsev7 reports much higher values of a t

, of the order of 10* 2, for copper at

Fig. 5. Creep rates versus applied stress for copper at 76 K; effects of

time interval on rates are illustrated. Data of Yen 1,3 also

included.

1 180

56
8800

to

<x

D5

8700

8600

8500

8400-

8300t

300 K

«. = 42 4 MP*

8200

1000 10000

Time, min

50000

Fig. 6. Dependence of creep scrain on Log cime for copper under a stress of

42.4 MPa at 300 K.

temperatures ranging from 1.5 to 4.2 K. The higher values led to the

suggestion that dislocation tunneling played a role in low- temperature creep

processes

.

7

If 4-K creep scrain is plotted linearly versus time, the curves of 4-K

measurements are significantly different from chose at higher temperatures.

A very brief transient period is noticed. The length of this transient

period was less chan 500 min in all tests. Owing to the diminishing creep

races with cime, creep strains did not exceed 150 x 10* e following the

instantaneous scrain, < 0.

Creep Behavior - 76 K

Creep scrain is linearly proportional to the logarithm of the cime at

76 K within experimental uncertainties. Evidence for this dependence is

shown in Fig. 4. The coefficient, a! , for the logarithmic dependence is

plotted versus applied stress in Fig. 3. While a t at 76 K has much the same dependence on stress as a t at 4 K, the magnitude of a t at 76 K is about

3 x 10' 5 larger chan at 4 K. The lowest stress for the Wyatt data 4 is about

62 MPa; therefore, these data are not included in this figure. However, a t is also lower (- 3 x 10’ 5

) chan our data for high stress levels.

When creep scrain at 76 K is plotted linearly versus Cime a larger

transient region is apparent, compared to curves at 4 K . It appears chat

there is a linear region, corresponding to a steady-state creep race, in

each plot of scrain vs. time. Yet, as illustrated in Fig. 5. careful data

analyses yield ever -decreasing creep races. In Fig. 5 the data of Yen ec

al. 2,3 obtained over the cime interval from 3000 to 12 000 min. are plotted

and joined by a curve. Our creep-rate data, obtained over the same cime

interval from our curves, are plotted for these specimens. They conform

excellently to the Yen data. But, at Longer periods of elapsed cime. lower

values of the creep race are calculated and included in Fig. 5. Creep races

as low as 1 x 10' 11 s' 1 for stress levels above 33 MPa at 76 K have been

obtained. Therefore, we conclude chat there is no steady-state creep stage

in copper at the stress levels used in this study. At these stress levels,

the creep race is continuing to decrease as evidenced by the linear dependence of scrain on log cime (Fig. 4).

1181

57
Creep Behavior: Room Temperature

A typical creep curve for room temperature Is shown in Fig. 6. At

ambient temperatures, the curve of creep strain versus time can be portrayed

as logarithmic at times less than approximately 10 4 min. A plot of a t obtained from Eq. (2) (Table 1) analyses versus applied stress is shown in

Fig. 3. The slopes of a t versus oa (see Fig. 3) gradually decrease with

decreasing temperature, and the absolute values of a t increase with increasing temperature. However, at elapsed times greater than about 10 4 min, the dependence of strain on time increases, relative to a logarithmic

dependence. This, again, provides an opportunity for the assessment of

steady-state conditions. At 300 K, the length of the transient region is

longer than that of 76 K curves. Our calculated values of creep races range

from 2.5 x 10* 11 to 6 x 10’ 11 s' 1 over a stress range from 20 to 45 MPa at

300 K for time intervals from 20 to 50 x 10 3 min.

The extrapolated value of <? a at (e 0 ) u* - 0 may be regarded as the

elastic limit for high strain rates (o 0 ). At room temperature a 0 is nearly

zero; at 76 and 4 K a Q is nearly equivalent (within data scatter), about

17 . 5 MPa.

Gohn and Fox 9 and Davis 10 have conducted room temperature creep tests

of copper for longer periods of time (15 x 10 s and 6 x 10 4 min respectively) . The strain versus time curves of Gohn and Fox for room temperature

creep tests exhibit decidedly decreasing strain races with time. No steadystate range is observed. However, if one uses the tangent of the curve at

12 x 10 s rain for their data, a strain race of 0.25 x 10* 11 s* 1 is calculated

for an applied stress of 69 MPa at room temperature. Similar low creep

races may be obcained from the data of Davis at lower stress. Thus, longer

periods of time lead to diminishing creep rates; steady-state conditions are not achieved at room temperature.

Engineering Strain and Creep Conditions

Finally, the strains one may expect at these three temperatures may be

summarized for various stresses. At room temperature for a stress equivalent to the yield strength (30 MPa), the maximum total creep strain for 20

years service is approximately 0.02 (including « Q - 0.004) from our data; if

the lower creep races from longer- time data 9,10 are used, a total creep

strain of less chan 0.006 (including « Q - 0.004) is estimated. At stresses

of 1.5 a

y our races approximately doubled; thus, a maximum extrapolated (20

years) strain would be 0.04; using the lower rate data®

•

10 this strain would

be reduced to 0.01.

At 76 K che creep races are similar to room temperature, therefore the

total predicted strains would reflect che difference in instantaneous strain

(about 0.0035 at 76 K) . At 4 K the creep races are less but the instantaneous strain is equivalent to that observed at 76 K. Total strain for a stress level of che yield strength (37 MPa) for 20 year service would be

predicted to reflect only che instantaneous strain (- 0.0035), since the

creep race exhausts relatively quickly at 4 K.

Ue used Eq. (2) (Table 1) for the analyses. At 76 K and room temperature, che contribution from the a s term in Eq . (2) was small compared to

that from che a t term for che typical test duration. Before 10 000 min, che

contribution from the a s terra is generally negligible, and between 10 000

and 30 000 min, the increase in creep strain due to che a s term is usually

no more chan about one -half of che increase due to che a t term. Beyond

about 40 000 min, however, che a s term becomes dominant, in part because log

t increases very slowly at chat point compared to c. Because a, is smaller

at 4 K , the a s term is important after shorter periods of elapsed time than

at 76 K and room temperature.

1 182

58
SUMMARY

The creep characteristics of 99.99% pure, oxygen- free copper (C10400) were studied ac 4 and 76 K and ac room cemperacure ac stress levels between

0.7 and 1.5 of the yield strengths. Major conclusions are:

1. At all temperatures the primary stage is best described by a logarithmic dependence of creep strain on time.

2. No steady- state creep was detected at any temperature. Creep rates

continued to decrease or to exhaust at each temperature as a function of time.

3. At 4 K logarithmic creep was followed by a relatively quick exhaustion of creep. At 76 K logarithmic creep continued for the duration

of the tests, within the limits of experimental imprecision. At

room cemperacure creep races increased from the logarithmic creep of

the primary stage, but did not attain steady-state conditions.

4. Total creep strains (including strain on application of the load) of

less than 0.020 are predicted for o a - ay ac room cemperacure for a service time of 20 years. At 76 K a total creep strain of less chan

0.019 is predicted; at 4 K, the total creep strain is less chan

0.004 when a a - Oy.

ACKNOWLEDGMENTS

This project was supported by the Office of Fusion Energy (DoE).

Elizabeth S Drexler assisted excellently in data analyses.

REFERENCES

1. J.K. Tien and C.T. Yen, "Cryogenic Creep of Copper," Adv. Cry. Eng.-

Maters. 30 : 319-338 (1984).

2. C. Yen, T. Caulfield, J.K. Tien, L.D. Roth, and J.M. Wells, "Cryogenic

Creep of Copper," Paper 8305-030, Mecals/Macerials Technology Series

American Society for Metals, Metals Park, OH (1983).

3. C.T. Yen. "Long-Term Creep of Copper at Cryogenic Temperatures," Ph.D.

Thesis. Columbia U., New York (1983).

4. O.H. Wyatt, "Transient Creep in Pure Metals," Proc. Phys . Soc. B66 : 459-480 (1953).

5. C. Yen, T. Caulfield, L.D. Roth, J.M. Wells, and J.K. Tien, "Creep of

Copper at Cryogenic Temperatures," Cryogenics 24:371-377 (1984).

6. V. A. Koval, V.P. Soldatov, "Jumplike Deformation of Copper and Aluminum

During Low-Temperature Creep," Adv. Cry. Eng. -Maters. 26:86-90

(1980).

7. V . I . Startsev, "Low Temperature Creep of Metals," Czech. J. Phys. B 31:

115-124 (1981).

8. R.W. Evans and B. Wilshire, Creep of Metals and Alloys . The Institute

of Metals. London (1985).

9. G.R. Cohn and A. Fox, "New Methods for Determining Stress -Relaxation ,

Matls. Res. Stds . : 957 - 966 ( 1961).

10.

E.A. Davis, "Creep and Relaxation of O.wgen-Free Copper," J. Appl.

Mech. 10 : A101 - 105 (1943) .

U83

59
.

-
STRAIN KATE EFFECT OR TEKSILE PROPERTIES

AT 4 K OP A VAMAS ROUND -ROB IN AUSTENITIC STEEL

R.P. Reed, R.P. Welsh and R.L. Toblar

Fraccura and Deformation Division

National Inscltuta of Standards and Technology

Boulder, Colorado USA

ABSTRACT

A high-strength austenitic steal with nominal composition Fe-25Cr- 14N1-

0.37N was included in an international round-robin measurement program.

Tensile and fracture toughness tests vere conducted at 4 K by leading low

temperature test laboratories in Japan, Europe, and the U.S. This paper

reports on the effect of strain rate on displacement-controlled tensile

tests at 4 K. Similar to other recent results on austenitic steels, a transition in ultimate strength was observed as a function of strain rate:

at higher strain rates the strength decreased about 10%. The onset of

discontinuous yielding was also strain-rate dependent. The reduction of

ultimate tensile strength is associated with the change from the nucleate to

film heat- transfer mechanism from the specimen surface to liquid helium.

INTRODUCTION

Low temperature test standards ara now in the process of being established both in the U.S. through the ASTM E-28 committee, in Japan through

JIS, and with active participation of European tast laboratories. One focal

group for discussion and round- robin testing leading to such standards is

the Versailles Advanced Material and Standards (VAMAS) organization. This

group selected a Japanese austenitic steel (Fe-25Cr- 14N1-0. 37N, YUS 170) as one of their two alloys for round- robin tensile and fracture toughness tests

at 4 K. 1 To complement these tests, we have extended our measurements at

NIST to study the effects of strain rate at 4 K since our previous research

identified a transition in ultimate strength as a function of strain rate

for austenitic steels. 3 Furthermore, many laboratories in the VAMAS round

robin conducted their tests at strain rates within the transition region.

This paper presents the NIST measurements and discusses them in terms of

round-robin test variability and future tensile- test standards.

EXPERIMENTAL CONDITIONS

The composition of the austenitic steel is Fe-25.06Cr-13.52Ni-0.76Mo0 . 78Mn- 1.12SI-0.015C-0.37N-0-035P-0.001S. Annealed specimens were received

from the National Research Institute for Metals (Tsukuba, Japan). The

Advances m C'votentc EngineerTM? iMaienalsi. Vol 36 1 06 1 EdueO t" R P Reed and F R. Fn-keii Plenum Pre>». Se» Vork. 1990

61
Fig. 1. Tensile yield strength at 4 K as a function of strain rate for Fe25Cr- 14Ni *0 . 37N austenitic steel.

average grain size was SO equiaxed grain structure vas observed in all

photomicrographs. Specimen diameter was 6.35 mm with a 37.5 mm gage length;

total specimen length was 87.5 mm.

Tensile tests were conducted in displacement control in boiling liquid

helium (4 K) using a screw-driven, commercial machine. The cryostat and

clamp-on. strain-gage extensometers have been described previously. 2 Elongation was measured over the entire specimen gage length. Reported

strain races refer to the nominal rate of plastic deformation within the

reduced section of the specimens.

EXPERIMENTAL RESULTS

Strain race has little effect on the tensile yield strength at 4 K

of the Fe-25Cr-14Ni-0. 37N alloy as shown for thirteen tests in Fig. 1. The

two data points of very low yield strengths (1200 and 1218 MPa) were closely

examined. Neither metallurgical variables (such as grain size), nor measurement variables (such as irregular stress-strain curve) can explain

these data. The most likely explanation is chemical inhomogeneicy within

the specimen gage lengths.

There is a transition in ultimate tensile strength as function of

strain race; at strain races in excess of about 1.3 x 10* 4 s* 1 the ultimate

strength begins to decrease (Fig. 2). The ultimate strength decreases about

10% at higher strain races and remains thereafter constant at strain races

in excess of about 4 x 10* 4 s* 1. Elongation, as shown in Fig. 3 increases at

higher strain races (> 4 x 10* 4 s* 1 ). This increase is not associated with

an increase in the reduction of area; reduction of area values vary from 31

to 48%, but show no trend with strain race.

Discontinuous yielding was observed during all tensile tests at 4 K. As described in an earlier paper, 3 the initiation of discontinuous yielding

in austenitic steels at 4 K is a function of strain race; both the initiation stress and strain exhibit transitions to lower values at higher strain

1062

62
2000

1700

t—i—i i r rr "i 1—r i i i i i |

Fe • 25Cf t4Ni o 3TN

I 1111 J 1 i 1111

10 10

Strain Rate, s

Fig. 2. ULcimate tensile strength at 4 K as a function of strain rate for

Fe-25Cr- 14Ni-0 . 37N austenitic steel.

rates. Engineering stress-strain curves at high and low strain rates are shown in Fig. 4. For the Fe-25Cr- l4Ni-0 . 37N alioy, the stress (o^) and

strain («^) at initiation of discontinuous yielding are plotted in Fig. 5a

and b, respectively. Similar to the other alloys, both initiation stress

and strain rapidly decrease as the strain rate increases above 10* 4 s* 1

.

Fig. 3. Tensile elongation as a function of strain rate for Fe-25Cr-14Ni

0.37N austenitic steel.

1063

63
Fig. 4. Tensile stress - strain curves (engineering) of Fe - 25Cr- 14Ni -0 . 37N

austenitic steel at tvo strain rates at 4 K.

DISCUSSION

Alloy Fe-25Cr- 14Ni-0 . 37N is stronger than most ocher austenitic alloys

chat we have studied. It exhibits a transition of ultimate strength (o M

) to

lower values at higher strain rates, similar to other austenitic steels.

However, unlike ocher austenitic alloys, the transition of the ultimate

strength of the Fe-25Cr- 14N1-0. 37N alloy occurs at lower strain rates. The

range of the strain rate in the transition region (to lover values of au ) encompasses the strain races used by most laboratories in the VAMAS roundrobin tests of this alloy. Therefore, the variability of strain rates from

contributing laboratories is very likely to account for much of the data

variability found in the round- robin tests. Ve discuss the ramifications of

this and present our interpretation of the divergent tensile behavior at 4 K.

Discontinuous Yielding

The sudden surge of localized catastrophic slip during tensile tests at

4 K of most structural alloys produces local temperature Increases of the

order of 50 K. At higher strain races, the nature of the temperature

increases change from spike -like to wave -like; chat is, the strain duration

of each temperature increase arising from discontinuous yields lengthens.

Reed and Walsh 3 and Ogata, et al 4 report these characteristics.

As shown in Fig. 6, the stress and the strain chat are associated with

the onset of discontinuous yielding follow trends very similar to those of

304L, 310, and 316LN. Serrated yielding in all alloys initiates at lower

values of a ^ and for strain races higher chan about 10" 4 s* 1

. At lower

strain races there is no detectable effect on initiation; at strain races higher chan about 4 x 10’ 4 s" 1

, minimum values of and are maintained.

Therefore, a transition from higher to lower initiation stress and strain is

observed over the strain rate range from 10

* 4 to 4 x 10* 4 s" 1 for all

austenitic steels that we have measured.

1064

64
2000

1300

1—i—r-| i i i • ~ • i r—i—r-r I' M

i i i i i i 1 10'

J L

Strain Rate. S'

1 10

Fig. 5. The scress (a) and strain (b) corresponding to initiation of discontinuous yielding at 4 K as a function of strain rate for alloy

Fe-25Cr-14Ni-0.37N.

We discuss the mechanism of discontinuous yielding in another paper of

*.his volume. 6 The onset of discontinuous yielding is attributed to thermal

instability; this produces mechanical instability resulting from localized

overheating of slip bands.

Reduction of Uitimate Strength

In Fig. 7 the ultimate tensile strength at 4 K is plotted versus

strain rate for the Fe-25Cr- lANi-0 . 37N alloy and austenitic alloys that

1065

65
were previously studied*. The chree lower -strength elloys (304, 310, end

316LN) oeintained high ultimate strengths until the strain rate exceeded

2 x 10' 3 s

1

. Thus, in the current draft of the 4-K tensile test standard

proposed to the ASTM committee E-28 on low- temperature test standards, it is

recommended chat strain-rates do not exceed 10’ 3 s* 1

.

The reduction of ultimate tensile strength in the higher strength Fe2SCr- 14N1-0. 37N alloy is associated with discontinuous yielding. At strain

rates below 10' *s~ l serrated yields were rarely observed; occasionally one serrated yield occurred and initiated specimen failure. Therefore, based on previous studies 3 * 6 ,' we can conclude chat at lower strain races the speci- men cemperature was maintained at (or very near) 4 K. At higher strain

races discontinuous yielding was always observed, and deviations of flow

stress to lower values accompanied them. This distinction between lack of

serrations and higher flow stress on the ocher hand, and serrations and

lower flow stress on the other is illustrated by the stress-strain curves of

Fig. 4.

In lower-strength austenitic steels, the serrations are smaller (25-

150 MPa as opposed to 300-350 MPa for this alloy). In the lower -strength

alloys there is some plastic deformation with attendant strain hardening in

the interval between elastic reloading (following a serration) and the

initiation of a new serration. These strain-hardening regions result in

similar stress -strain curves for both high and low strain rates.

Fig. 6. The stress and strain corresponding to initiation of discontinuous

yielding at 4 K as a function of strain race, comparing alloy Fe25Cr- 14N1-0. 37N to other austenitic alloys (304L, 310, and 316L21)

.

1066

66
Fig. 7. Ultimate tensile strength at 4 K as a function of strain rate for

selected austenitic steels.

The large magnitude of the serrations in the present alloy strongly

suggests very high temperature spikes. For alloy 304L2

. load drops of about

50 MPa produced localized temperature spikes of about 50 K; load drops of

about 150 MPa produced temperature increases of nearly 100 K. This suggests

that local temperature increases of over 150 K may result from discontinuous

yie lding in the Fe-25Cr-14Ni-0. 37N alloy. Therefore, since the ultimate

strength of this alloy is approximately 1400 MPa at 76 K, it is not sur- prising that the failure strength is reduced under discontinuous yielding

conditions at 4 K. The increase of tensile elongation with increasing

strain rate (Fig. 3) is also expected to originate from the large temperature increases associated with discontinuous yielding.

Effects on Round-Robin Interpretation and Standards

The range of strain rates used by the 16 laboratories participating in

the VAMAS round-robin 1 of tensile tests at 4 K was 1.3 to 11.1 x lO’^s* 1

. It is evident from Figs. 2-6 that the data from all laboratories should not

be expected to coincide, but should be expected to be dependent on the

strain rate. For example, the ultimate strengths at 4 K of alloy Fe-25Cr14N1-0.37N reported by the 16 laboratories ranged from 1791 to 1960 MPa. Of

the 42 data points for this property, only three exceeded 1900 MPa (1920,

1920, and 1962 MPa). In contrast, at low strain rates (< lO'^s' 1 ) all of

our measured values exceeded 1900 MPa. Therefore, the round- robin test

program on the Fe-25Cr- 14Nt-0 . 37N is inconclusive for assessment of the

tensile data associated with specimen failure. The reported data are biased

toward lower (approximately 10%) values of ultimate strength and toward

higher values (approximately 25%) of elongation.

The results reported in this paper indicate that the tensile test

procedures at 4 K commonly used in many laboratories need to be revised.

Many laboratories begin testing at 4 K at low strain rates, then increase

the strain rate by an order of magnitude following measurement of the yield

strength. This increase of strain rate is performed to conserve liquid

helium and test time. But either the exact opposite of this procedure or continuous maintenance of a single slow strain rate (<10* 4 s* 1 ) is necessary

to avoid serious specimen thermal imbalances. The thermal Imbalances that occur in specimens during testing of high-strength austenitic steels have

1067

67
been shown in chis scudy co r«sult in lower ultimate strengths and higher

elongations. The currently proposed draft standards for low temperature

resting now suggest strain rates below 10* s s* 1

; this should be changed to

10* 4 s* 1 for high strength steels.

Two challenges still remain for future research on low temperature

tenjile standards: (1) to identify, quantitatively, the conditions under,

which lower strain rates are necessary during tensile tests at 4 K, and

(2) to extend the strain- rate studies to nonferrous materials, particularly

high-strength, low -conductivity alloys (such as Tl-based alloys).

CONCLUSIONS

In the very high-strength Fe-25Cr- 14NI-0. 37N austenitic alloy used in

the VAMAS round-robin tensile test program, reduced strength and Increased

elongation were measured at strain rates an order of magnitude lower than

usual. Two major conclusions result from this scudy:

1. The variability of the VAMAS round-robin tensile data (other chan

yield strength) is partially attributable to the different strain

races used by each laboratory. Therefore, these results are inconclusive with respect to analyzing material and inter laboratory test

r ep roduc ib i 1 i ty

.

2. The tensile test standards currently under development worldwide

need to be revised to reflect the effects of strain rates above

lO'^s* 1 in high-strength, ductile austenitic steels.

ACKNOWLEDGEMENTS

The authors thank K. Ishikawa, K. Nagai, and T. Ogata of Japan's

National Research Institute of Metals for supplying tensile specimens and

mutual cooperation on chis program.

REFERENCES

1. VAMAS Superconducting and Structural Materials, The Third TVP Meeting,

National Research Institute for Metals, Science and Technology

Agency, Japan (May 30, 1988).

2 . R.P. Reed and R.P. Walsh, "Tensile Strain-Rate Effects in Liquid

Helium," Adv. Cryo. Eng. -Mater. 34:199-208 (1988).

3. R.P. Reed and N.J. Simon, "Discontinuous Yielding in Austenitic Steels

at Low Temperatures," Cryogenic Materials '88, Vol. 2, International

Cryogenic Materials Conference, Boulder, CO (1988), pp. 851-863.

4 . T. Ogata, K. Ishikawa, 0. Umezawa, and T. Yuri, "Effects of Specimen

Geometry on Temperature and Discontinuous Deformation during Tensile

Tests at Liquid Helium Temperature," Adv. Cryo. Eng. - Mater. 34:209-

215 (1988).

5. R.P. Reed and N.J. Simon, "Discontinuous Yielding During Tensile Tests

at Low Temperatures," Adv. Cryo. Eng. -Mater. 36: (1990).

6. D.T. Read and R.P. Reed, "Heating Effects during Tensile Tests of AISI

304L Stainless Steels at 4 K," Adv. Cryo. Eng. -Mater. 26:91-101

(1980)

.

1068

68
LOAD -CONTROLLED TENSILE TESTS OF AUSTENITIC STEELS AT 4 K

H.M. Lee

Korea Standards Research Institute

Taejon, Korea

R.P. Reed

National Institute of Standards and Technology

Boulder, Colorado

J.K. Han

Research Institute of Industrial Science and Technology

Pohang, Korea

ABSTPACT

Load-controlled tensile tests were conducted at 4 K on high-strength

austenitic steels. The race of loading was varied from 5 to 5000 N/s.

This change of loading race affected the onset of discontinuous yielding

and, in turn, the fracture characteristics of the steels. Ultimate strength

decreased at higher loading rates. The role of discontinuous yielding in

affecting the dependence of these properties on loading rate is discussed.

If conventional pressure vessel codes are used for the designation of

structural design stresses, these data suggest that the ultimate strength,

not the yield strength, controls this assignment.

INTRODUCTION

To simulate the operating conditions in high- field superconducting

magnets, load-controlled tensile tests were conducted on selected strong

austenitic steels. Most tensile tests at cryogenic temperatures have been

carried out under constant displacement rate. In many practical applications such as superconducting magnets and pressure vessels, however, the

structural material experiences unrestricted force. It is, therefore,

desirable to characterize the load-controlled tensile properties at cryogenic temperatures. Earlier research 1 on lower- strength austenitic stainless steels (304L, 310, 316LN) indicates that, at higher load rates, the

ultimate tensile strength and elongation are affected; the ultimate strength

is reduced, and the elongation increases Tensile yield strength and

reduction of area are relatively unaffected at 4 K by load rate. The effect

of load rate on initiation of discontinuous yielding was documented and

associated with the reduction of ultimate strength.

Advances in C'votenu Enfineermt iWaienatsi. \ ol 36 Edited b> R P Reed and F R Fickett Plenum Press. Ne» York. 1990 1273
Table 1. Material Characterization

Alloy

Cr Ni Mn Mo

Alloy Concent,

N C S wt.%

P Si

Crain

Size

Other (jxa) Hardhess

(Rfi)

Fe-18Cr3Ni-13Mn

18.1 3.3 13.2 0.1 0.37 0.04 0.005 0.028 0.52 0.18 No, 48

0.15 V

93

Fe-21Cr12Ni-5Mn 21.2 12.4 5.0 2.2 0.31 0.04 0.015 0.026 0.49 34 98

Fe-20Cr25Ni-6Mo 20.3 24.7 1.6 6.3 0.19* 0.02 0.002 0.021 0.43 0.20 Cu 45 88

a; independent N analyses yielded 0 198, an average of three data secs.

There is concern chac. under constant load-rate conditions, the initiation of a discontinuous yield could lead directly to catastrophic fracture

at 4 K. A single discontinuous yield originates locally along the reduced

section of the specimen, similar to a shear band. The deformed area throughout the specimen cross section had a band width (of the temperature

rise) about 10 mm 2

. Local temperature rises within this thermal band have

been measured to be between 50 and 150 K. 2

* 4 In higher- strength alloys and

at higher stresses in lower yield-strength alloys, the magnitude of the load

drops and corresponding temperature spikes are greater. The higher strength

of the austenitic alloys is normally achieved by nitrogen addition. Therefore, the flow and ultimate strengths of these alloys are expected to be

strongly temperature dependent. Stronger temperature dependence of the flow

strengths, coupled with larger load drops and temperature spikes, is ex- pected to produce a greater dependence on strain or load rate during 4 K

tensile tests. This paper explores this dependence for higher strength

austenitic alloys.

EXPERIMENTAL PROCEDURES

The Fe-20Cr-25Ni-6Mo alloy was supplied in 38-mm-thick plate from a commercial heat. Its chemistry (Table 1) conforms to UNS N08367. All

alloys tested had been annealed condition; details of heat treatments have

been previously reported. 5

Round, tensile specimens were machined to a 6.35-mm diameter along a gage length of 41.9 mm. The entire specimen length was 71.4 mm.

The cryostat equipment has been described elsewhere

.

1 * 3 • • Tests were conducted in boiling liquid helium and liquid nitrogen. Specimen strain was measured with clip-on, strain-gage extensometers : one extensometer with a

2. 5 -mm span and a gage length of 25.4 mm was used to measure yield strength.

The second, high- range extensometer, with a span of 17.5 mm and a gage

length of 38 ram, was used to measure specimen strain throughout the tensile

test. Elongation was measured over a specimen-gage length of 3.8 cm. Reported strain rates refer to the nominal rate of plastic deformation

within the reduced section of the specimens and can be measured directly in

load-controlled tests.

A programmable servohydraulic testing machine (maximum load 250 kN) was

used. With a high - resolution function generator, loading races were varied

from 0.5 to 5000 N/s . The machine was programmed for a ramp

1274

70
80

60

40

20

0

Fig. 1. Tensile properties at 4

and 76 K versus load rate.

Fig. 2. Tensile properties at 4

and 76 K versus Load rate.

mm/s

Loading Rate (N/s)

Fig. 3. Tensile properties at 4 and 76 K as a function of load rate.

1275

71

Reduction

ol

Aiea

Elongation

(%)
function in a Load-control mode; that is, tha machine uniformly incraasad

tha Load on the speciman until failura. Tha maximum actuator velocity is

asti -maced to be 50 mm/s . Tha strain rata in tha Linaar-alastic region was

8 x 10** s* 1 co 8 x 10* 4 s' 1

, corresponding to loading rates of 0.5 N/s to

5000 N/s.

Load and strain ware recorded with x-y recorders. During loadcontrolled tescing, che strain rate was measured just before discontinuous

deformaclon. A digital oscilloscope or strip-chart recorder was used for measuremenc of scrain rate.

EXPERIMENTAL RESULTS

The effects of load race on che censile yield screngch, ultimate

screngch, reduction of area (R.A.), and elongation are summarized in

Figs. 1-3. All alloys were tested ac both 4 and 76 K. Ac boch temperature*

che yield strength is independent of load rate; che ultimate screngch

decreases with increasing load rate. In one case (alloy 20Cr—5Ni—6Mo

,

Fig. 3) at higher load races, che ultimate strength at 76 K is higher than

ac 4 K.

With increasing load rate che reduction of area tends to increase ver

slightly, except for che large increase measured for alloy 18Cr-3Ni-13Mn ac

76 K and che highest load race (5000 N/s). Elongation decreases with

increasing load race at 76 K, but remains constant or increases (Figs. 3, 4 with increasing load race at 4 K.

The tensile data for displacement-controlled tests ac 0.008 mm/s are provided in Figs. 1-3. The ultimate censile strength at this displacement

rate for 295 , 76 , and 4 K are shown in Fig. 4. The more ductile 20Cr-25Ni6Mo alloy has lower censile strength ac low temperatures, but all alloys

have almost equivalent strength at room temperature.

Fig. 4. Ultimate censile strength (displacement control) vs temperature.

1 276

72
a

b

Fig. 5. Engineering stress * strain curves at 4 and 76 K for various rates

of loading (a) alloy 18Cr-3Ni-13Mn, (b) alloy 21Cr-12Ni-5Mn.

Stress-strain curves (engineering) are shown for various load races in

Figs. 5a—b. At 4 K, alloys 21Cr-l2Ni-5Mn and 18Cr-3Ni-13Mo in displacement

control exhibit many discontinuous yields. In load-controlled tests at 4 K,

all specimens fractured during the first discontinuous yield. No discontinuous yielding was detected in any alloy at 76 K.

The occurrence of discontinuous yielding was sometimes ambiguous.

Therefore, a digital oscilloscope was used to verify discontinuous yielding.

Figure 6 shows load and strain signals versus time, obtained from the alloy

of Fe-18Cr-3Ni-13Mn. At lower load races there is a clear discontinuity of

both load and strain with time; at the highest rate (5000 N/S) no discontinuity was observed, yec the load decreased. This indicates that the

1277

73
Fig. 6. Load (curves) and strain (lower curves) versus time for alloy

18Cr-3Ni-13Mn during pLastic deformation at 4 K for various load

rates (left: 5 N/s; center: 500 N/s; right: 5000 N/S).

actuator velocity (50 mm/s) was not large enough to keep pace with the

increased strain rate of about 7 x 10* l s' 1

.

Maximum internal specimen temperatures during testing at 4 K for the

three alloys were estimated by comparing the magnitude of the load drop with

the temperature dependence of the flow strength. The results are shown in

Fig. 7. Both the magnitude of the load drops and the estimated internal

specimen temperature increase (almost linearly) with increasing loading race

for all alloys.

The initiation of discontinuous yielding is dependent on the strain

race of plastic deformation. In Fig. 8 we plot this dependency for three

of the alloys plus alloys that had been previously tested at NIST 1 in load

control. All have transitions from higher initiation stress and strain to

lower values at about 10'*s* 1

.

Fig. 7. Estimated maximum cemperature of specimens during deformation at

4 K as a function of race of Loading.

1 278

74
a

Engineering Strain Rate iust Detore the First Serration is* 1 ' b

Fig. 8. The scress (a) and scrain (b) co initiate discontinuous yielding

for selected alloys versus specimen scrain rate at initiation.

DISCUSSION

Discontinuous Yielding and Yield Strength

In displacement -control studies, the scress (?i) to initiate discontinuous yielding is proportional co the yield strength (oy) . In particular,

we reported6 that

oi - o

y + 320 MPa (1)

for over 80 measurements of austenitic steels using low scrain rates.

1279

75
Generally, the flow strength (o) is considered a sun of contribution* from

strain hardening (o^) and solution hardening (of)

o - <? d + a f. (2)

The solution-hardening contribution may be equated to the yield strength

(of - ay ) . Then, when a - a^ at the initiation of discontinuous yielding,

odis a constant, equal to 320 MPa.

The yield strength at lower load rates for the austenitic alloys of

chis study plus a previous study at our laboratories 1 is plotted in Fig. 9 versus che initiation stress (c^). A linear relation is obvious, with the

intercept of 350 MPa at Oy - 0. Within experimental uncertainty, the

results from the previous study* and from these data are identical. Therefore, the strain hardening contribution to initiate discontinuous yielding

is constant, independent of strength level and of metallurgical variables,

such as stacking fault energy and grain size. However, external variables,

such as strain rate, heat transfer coefficients, and test temperature do

affect and o d

. 1 * 4,4

Discontinuous Yleldlne and Strain Rate

At higher load races, che stress and strain (<^> to initiate discontinuous yielding declines. This corresponds to che influence of strain race

in displacement control tests. 3,4 Under displacement control, a transition

from higher to lower values of aoccurs at strain races of che order of

10'*s* 1. If che strain race is measured at a ^ and plotted versus a

i

(Fig. 8a) and (Fig. 8b) for load-control che data, a similar transition

is apparent. In previous work, 4 we have suggested chat the transition may

be caused by a change of local heat transfer in boiling liquid helium from

nucleate to film boiling conditions at che microscopic level (at slip bands

within grains). This was argued because che macroscopic work and heat

balance at 4 K (considering work performed, stored energy, and solid and

liquid heat conduction) does not produce sufficient heating to account for a transition in che heat transfer process. This process is also discussed in

another paper of these proceedings. 7

Discontinuous Yielding and Ultimate Strength

The major distinction between displacement and load-controlled tests

at 4 K is che effect on tensile fracture characteristics . In displacementcontrolled tensile tests, che sudden strain associated with a discontinuous

yield results in a sudden reduction of che applied load. That is, che load

train experiences an effective relaxation, dependent on its effective stiffness. This permits time for che local deformation band within che specimen

to achieve thermal and mechanical equilibrium at a higher temperature and

lower load level. Equilibrium is realized when che local strength at che

higher temperature is balanced by che strain hardening of che deformation

band.

In load-controlled tests, following a discontinuous yield, che machine

increases its race of displacement and attempts to maintain constant load.

This results in additional strain to che local shear band since its temperature is higher and effective stress level lower chan che surrounding

material. The additional local strain within the shear band produces

additional heating and reduction of effective strength. Thus, che process

may be cumulative and will usually result in fracture within che shear band.

Our results demonstrate chis effect; in load-controlled tests fracture

strength was usually associated with che stress necessary to initiate

discontinuous yielding.

1280

76
Figure 9. Stress to initiate discontinuous yielding versus yield strength

for alloys from this study and from Ogata et al. 1,4

Applications to En gineering Practice

The severe decline of the ultimate tensile strength (<7 U ) under load

control raises the issue of the application of ASME Boiler and Pressure

Vessel Codes, Section VIII. This code is sometimes used to ensure sound

design and material application practice in superconducting magnet construction.

The code usually requires operating stresses of either 1/2 to 2/3 Oy

or 1/3 «7 U. For all load rates, the ultimate strength ratio is lower than

the yie Id- strength ratio, in contrast to displacement-control results.

Therefore, if the Boiler and Pressure -Vessel Code is thought applicable for use in low temperature technology development, the ultimate strength con- trols alloy selection for load-controlled structural applications such as superconducting magnets. Ue suggest that two directions need to be taken to

resolve this dilemma: (I) tests on thicker section components are needed to

permit better development of thermal and mechanical scale-up models from

laboratory specimen data and (2) development of alternative or amended

design codes is necessary for superconducting magnet structures.

SUMMARY

To simulate operational conditions of a structural material in a superconducting magnet, load-controlled tests were conducted at 4 K on three

high-strength austenitic steels.

(1) The onset of discontinuous yielding was correlated to the yield

strength - <jy + 350 MPa) and to the strain rate; a

i

and

undergo a transition to lower values as the strain rate increases

above 10 s' 1

. (2) The ultimate strengths in load-controlled tests are associated with

the onset of discontinuous yielding. Failure occurs within the

deformation band of the original discontinuous yield.

1281

77
(3) The issue of the application of ASME Boiler and Pressure Vessel Code

to low temperature technology is raised by these data: under code

rules the ultimate strength, not the yield strength, controls the

assignment of design stresses.

ACKNOWLEDGMENTS

Ue thank John Grubb of Allegheny- Ludlum Steel for sending a place of

the Fe—20Cr-25Ni—6Mo alloy. Robert Walsh assisted and advised in the

measurement program. The Office of Fusion Energy partially sponsored this

program

.

REFERENCES

1. T. Ogata, K. Ishikawa, R.P. Reed, and R.P. Walsh, "Loading Rate

Effects on Discontinuous Deformation in Load-Control Tensile

Tests," Ad. Cryo. Eng. -Mater. 34:233-240 (1988).

2 . D.T. Read and R.P. Reed, "Heating Effects during Tensile Tests of AWI

304L Stainless Steel at 4 K, " Adv. Cryo. Eng. -Mater. 26:91-101

(1980).

3. R.P. Reed and R.P. Walsh, "Tensile Strain-Rate Effects in Liquid

Helium," Adv. Cryo. Eng. -Mater. 34:199-208 (1988).

4. T. Ogata, K. Ishikawa, 0. Umezawa, and T. Yuri, "Effects of Specimen

Geometry on Temperature and Discontinuous Deformation during

Tensile Tests at Liquid Helium Temperature," Adv. Cryo. Eng.-

Mater. 34:209-215 (1988).

5 . D.T. Read and R.P. Reed, "Toughness, Fatigue Crack Growth, and Tensile

Properties of Three Nitrogen-Strengthened Stainless Steels at

Cryogenic Temperatures", The Metal Science of Stainless Steels . Metall. Soc. AIME, New York (1979), pp. 92-121.

6. R.P. Reed and N.J. Simon, "Discontinuous Yielding in Austenitic Steels

at Low Temperatures," Cryogenic Materials '88, (Shenyang, China),

International Cryogenic Materials Conference, Boulder, CO (1988),

pp. 851-863.

7. R.P. Reed and N.J. Simon, "Discontinuous Yielding during Tensile Tests

at Low Temperatures," Adv. Cryo. Eng. -Mater. 36: (1990).

8. "Basis for Establishing Allowable Stress Values," ASME Boiler and

Pressure Vessel Code, Section VIII, Division 1, Appendix P,

American Society of Mechanical Engineers, New York.

1 282

78
EFFECT OF PROCESSING ON 4-K MECHANICAL PROPERTIES OF A

MICROALLOYED AUSTENITIC STAINLESS STEEL'*’

P.T. Purtscher

Advanced Sceel Processes and Produces Research Center

Colorado School of Mines

Golden CO

M.C. Mataya

Rockwell International

Rocky Flats Plant

Golden, CO

L.M. Ma* and R.P. Reed

National Institute of Standards and Technology

Boulder CO

ABSTRACT

Tensile and fracture toughness tests were performed at 4 K as a function

of hot-rolling temperature and compared with similar data for the steel in

the annealed condition. Results show that the properties in hot* rolled

conditions were comparable to annealed steel so there was no advantage to

annealing after processing. Microalloy precipitates were effective at

controlling grain size during processing at temperatures below 1150 C, and

did influence the flow stress at 4 K.

INTRODUCTION

Development of austenitic stainless steels for superconducting magnet

cases has concentrated on the effect of alloy composition on mechanical

properties in the annealed condition after processing. The familiar NIST

trend line for fracture toughness as a function of yield strength at 4 K was generated for a series of annealed austenitic stainless steels, AISI 304-

type with different interstitial contents [1].

Limited data on the effect of thermoo-echanical processing of austenitic

steels are available. Ogawa and Morris [2] have studied the influence of

processing on the 4-K mechanical properties of a high -manganese , austenitic

stainless sceel. They found that reducing the hot-rolling temperature from

1250 to 1150*C refined the grain size, producing an increased yield

strength. However, fracture toughness decreased rapidly. Room temperature

deformation (30%) improved the 4-K strength- toughness relationship.

contribution of NIST, not subject to copyright

on leave from Institute of Metal Research Academia Sinica, Shenyang,

P.R. China

Advances m Cryogenic Engineering /Materials). Vol J6 1291

Edited by R P Reed and F. R. Fickett Plenum Preai. Ne* York, 1990

79
Table 1. Tensile Properties of As-Received Bar ac 4 K

Location 0 . 2%Y . S . U.T.S. Elong. % R.a. % (MPa) (MPa)

#1 Outer diameter

#2 Mid- radius

#3 Center

»4 Mid-radius

«5 Outer diaaeter

1540

1212

1107

1171

1503

1877

1810

1794

1792

1858

24.4

16.3

13.3

20.8

28.4

33.7

30.4

26.6

37.1

40.4

This investigation will determine the effect of hot-rolling conditions

on the mechanical properties of an austenitic, structural steel for 4-K

service. The starting temperatures for rolling (1150 to 950*C) are lower

than that used by Ogawa and Morris [2]. If the mechanical properties after

hot rolling are comparable to the properties in the annealed condition, then

it should be possible to save the time and expense of a final annealing

treatment before these types of steels are fabricated into structures.

MATERIALS AND PROCEDURES

Starting Material : The steel, Fe - 22Cr - 1 3Nt - 5Mn- 2 . 3Mo -0 . 2 INb-O . 1 8V- 0 . 03C

0.31N, was provided by a commercial, specialty- steel maker in the form of a 76 -mm-diameter bar in the mill-annealed condition. The high concentrations

of V and Nb control the grain size during processing and contribute slightly

to the room temperature yield strength. The as-received microstructure of

the bar was 100 % austenitic at the outer diameter, but the center of the bar

contained 0.1 to 0.2 volume fraction of sigma phase in the austenitic

matrix. A hardness profile through the cross section of the as-received

material showed a higher hardness, RA 61, at the outer surface decreasing to

Ra 54.5 in the center where the sigma phase was present. Table 1 shows the

4-K tensile properties, measured as a function of location through the

diameter of the as-received bar. The large variation in tensile properties

for the bar at 4 K could cause a problem when the steel is fabricated for

cryogenic service.

Material Processing : In the first step, the round bar was heated to 1200*C

and flattened in two hits by a steam hammer to a thickness of 38 mm. Next,

Fig 1. Microstructure of steel after hot rolling from 1050 C and water

quenching. Small remanent of sigma phase (arrow) and transgranular

precipitates are visible in austenitic matrix, etch: HNO 3 +H 2 O (3 8 volts

1 292

80
sections were reheated to 1150, 1050, or 950*C for 30 min before final hot

rolling in two passes to the final thickness of 28 mm. Immediately after

rolling, the place was straightened in the steam hammer. The hoc rolling

and straightening typically took about 2.5 min before the section was quenched into water. The finish temperatures before quenching were 875,

790. and 775*C respectively. One section rolled from 1150*C was annealed

after processing for comparison to the hoc-rolled material and is designated

condition A. The sections rolled from 1150, 1050, and 950*C are called

conditions D, B, and C, respectively.

Test Procedures : Two round tensile specimens oriented across the rolling

direction and three 12 . 7 -ram- chick compact specimens (C-T) with T-L

orientations were taken from the cop half of the plate. The »1 and *3 C-Ts were taken from the edges of the plate, and #2 came from the center of the

width. Additional C-Ts were machined from edge locations of selected

places. Longitudinal tensile specimens were machined from the different

conditions so there are five specimens equally spaced across the width of

the bottom half of the places, *1 and *5 at the edges of the place and »3 in

the center.

The toughness tests were conducted according to ASTM E813-81 with the

unloading compliance method for measuring crack growth during the test. The

only problem in testing was chat the crack resistance curve usually

contained only three or four data points between the exclusion limits of

0.15 and 1 . 5 ram sec by the standard. For a valid test, the standard

requires at least four data points between the exclusion limits. Jj c was defined as the intersection of straight line drawn through the data within

the exclusion limits and the blunting line as defined in the standard. More

details of the tensile and fracture toughness testing are described in a separate paper in this volume [3].

RESULTS

Material Characterization : The processed places had relatively uniform

microscructures with little evidence of sigma phase. Figure 1 shows a remanent of sigma after rolling from 1050*C and a random precipitate

distribution V and Nb(C,N) in the austenite chat was present in all the

plates. The grain size of the steel in all four conditions was between 40

and 60 um . The precipitates were effective at pinning the grain boundaries

and preventing grain growth during hot rolling and annealing. Annealing at

1 1 6

3

* C for 8 h was required to grow the grains significantly [4]. The

hardness of sections taken through a place from each test condition showed

that the materials were relatively uniform: for plate A -- RA59.5 to 62,

for place B -- RA 61 to 64.5, for place C -- RA 64.5 to 66, and for place D

-- Ra 59.5 to 63.

Tensile test : The results are summarized in the left half of Table 2 for

both longitudinal and transverse orientations. Hot rolling at lower

temperatures has increased the 0.2% FS with little change in ductility or ultimate strength at 4 K . Each of the places shows better uniformity than

the as-received bar. The tensile properties with the most variablicy are yield strength (0.2% FS) and percent elongation. Figure 2 shows the

variation about the average values in these properties in the longitudinal

orientation for each condition. Hoc rolling at lower temperatures reduced

the variations observed across the width of the plate.

Toughness test : The results are summarized in the right half of Table 2. The data for condition A (annealed) appear to be divided in two distinct

groups, one with an average Kj c of 110 MPaVm and a second with an average

Kj c of 144 MPa^m. There is no systematic variation in toughness across the

1293

31
Table 2. Mechanical Properties of Processed Places at 4 K

Tens l Le Prot»m.« Fraccure Touehness Proot r o i e s Spec . 0.2% FS LIS Clone

.

R . A. Spec . J Ic *Ic

**• 4 MPa e e fcJ/T 2 MPa-Zra

condition A -- annealed

AL- 1 1320 1795 28.8 31 . 7 A2

98 6 140.4 22.6

AL-2 1448 1784 22 .

4

42 .

5

A3

102.

2

143.0 13 6 AL- 3 1425 1799 26.4 46.2 A - 4 108 .

9

108 .

9

; 2 AL - - 1439 1764 23.0 39.7 A-5 58 .

4

108.1 6 . 1 AL * 5 1425 1779 23.4 51.6 A - 6 62 3 111.6 7 . 7 AT -2 1264 1 / 64 22 1 45.0 A - 7 58.4 107.6 5 3 AT - 3 14-1 1877 14 3 ‘*3 9 A * 8 109 .

5

148 0 13.4

A- 9 111.8 149 5 17.7

condicon B -- hoc rolled from : O o

e o 3L- 1 1714 1917 19.4 39 5 3- 1 •*8.6 98 6 BL-2 1704 1919 18.4 38.6 B - 2 *3 7 9 3 5 3 6 BL- 3 1687 1915 18 . 2 41 .

3

B- 3 *7.0 97 0 2.0

BL-4 1704 1913 12.9 44 .

4

BL- 3 1707 1912 18 . 2 42.3

BT-1 1677 1876 16.4 43.5

3T-2 1696 1867 16.2 *0.3

condition C

•• t10 ? ro Lied from 950*C

CL- 1 1800 1970 17.5 36.8 C- l 33 9 32.3 1 . 5 CL2

1807 1949 17.2 39.2 C - 2 3*. 2 82 . 7 3 2 CL- 3 1759 1904 19 8 35.4 C- 3 42.0 91 6 2.6

CL-4 1801 1952 13 0 34.7 C -4 32 . 5 80 6 2 . 1 CL- 5 1839 1967 17.1 42.7

CT- 1 1740 1968 9 6 39 . 1 CT- 2 1737 1971 15.3 35 . 1

cond; ton D DL1

•• rolled

1629

from 1 1914

130-C

16 7 47.6 D1

66 1 115 0 6.5

DL-2 1569 1894 21.0 *1.6 D - 2 72 5 120.4 9.2

DL- 3 1478 1816 19 5 -2 6 D- 3 SS 8 133.3 9.5

DL-4 1534 1888 15 . 3 45 3 D-4 61 . 2 110.6 5.0

DL- 5 1642 1902 12.1 40 4 D-5 59 8 109.4 5.3 DT1

DT-2

1574

1558

1829

1389

21 3 12.0

4 <* 3 *2 2

* Kt c (J) 2 - JI C x E where E - 200000 MPa

* T*“ (dJ/dalxt E/s Y

J > where sY - (0.2% FS l'TSl/2

width of the places. Figure 3 shows the toughness values plotted as a function of average, transverse 0.2% FS . The results from this study are compared with the original data used in the NIST trend line for AISI 304-

cype steels corrected for the same Jic* co ‘ K Ic conversion factor. The data

for condition A fit the previous trend for annealed steels, but the data for

conditions B, C, and D are above the extrapolation of the lower- strength

data

.

DISCUSSION

In a hot-rolled condition, austenitic steels should have higher yield

strengths chan the same steel after annealing. In general, fracture

toughness properties are inversely related to material strengths [5]. By

quenching directly after rolling, excess dislocations are retained in the room temperature structure and make additional dislocation motion more difficult, chat is, raise the flow stress compared to condition A. As the

1 294

82
LUo

SJ

So 1 O-o

cc C

1.5

1.4

1.3

1.2

1.1

1.0

0.9

0.8

0.7

0.6

0.5 As Received A 0 B C

Bar Plate Plate Plate Plate Annealed Annealed HR 1150 HR 1050 HR950

Fig. 2. The effect of processing on the variation in tensile properties, P (0.2% FS and % elongation), is shown as the ratio of the highest

and lowest value, Phigh and Plow- divided by P ave •

rolling temperature decreases, more dislocations should be retained in the room temperature structure and, as a result, the flow stress should increase

and fracture toughness would be expected to decrease. The annealing treatment should remove all excess dislocations that are produced in hot rolling.

The steel used in this study has an additional complication due to the pre- sence of V and Nb in the composition. These alloying element combine with C and N to form precipitates which stabilize the austenite, prevent Cr carbide

precipitation, resist grain growth, and strengthen the matrix.

If there were no Nb and V in the Fe- 22Cr- 13Ni -5Mn-2 . 3Mo-0 . 3N steel, the

0.2% FS could be predicted from the results of Simon and Reed [12] to be

1300 MPa with a S.D. of 40 MPa. The 0.2% FS in condition A was measured to

Fig. 3. Fracture toughness plotted against 0.2% FS Data for 304- type

steels are from ref. 1.

1295

83
250

Ifm

CLa

H

200 .

KX

Uia

z

>

100 .

<3

O

0 L_

900

o

0

4

4

§

i

r

4

t

Annealed 1

4

f

HOI

Rolled

4 Fe-18Mn-5Ni-l6Cr-0.22N

O Fe-22CM3Ni-5Mn-Q.3N

1 • i -- i 1000 1100 1200

PROCESSING TEMPERATURE, C

1300

Fig. 4. Quality index pLocted against processing temperature for either hot

rolling or annealing. Data for high-Mn steel is from ref. 2.

be as high as 1540 and 1448 MPa in the bar and plate, respectively. The

difference between the measured and predicted values can be attributed to

the Nb and V precipitates.

One effect of hot rolling was to make the tensile properties more uniform (Fig. 2). The increase in uniformity is related to the precipitation behavior of the steel. When the steel is annealed, there are fewer

nucleation sites for precipitates and the density of particles is probably

more variable. In the hot- rolled conditions, excess dislocations left from

processing can nucleate a more uniform distribution of particles.

The higher strength that comes from Lowering the rolling temperature

for this microalloyed steel does reduce J{ c , but the decrease is not as dramatic as Ogawa and Morris [2] observed. To help illustrate this point,

compare the quality index (QI) calculated for each test condition; see Fig.

4. The QI is the product of K^ c and 0.2% FS and has been used to relate the

influence that chemical composition and annealing temperature have on the

strength- toughness relationship in austenitic steels [3,6 ,7,8]. The data

from Ogawa and Morris [2] is included in Fig. 4 for comparison. Reducing

the rolling temperature from 1250 to 1150*C reduced the QI for Fe-18Mn-5Ni16Cr-0.22N (high-Mn steel) to a level below that of the annealed condition.

For the Fe-22Cr- 13Ni- 5Mn-0 . 3N steel in this study, the QI decreases slightly

as the rolling temperature goes down. Again, the larger variablity of the

data for condition A is obvious. There does not appear to be an inherently

better strength- toughness combination for annealed steel than for steel

processed in the 1150 to 950*C range. The overall higher QI for the high-Mn

steel compared to the steel in this study can be attributed to the different

chemical composition and different test orientation.

The scatter in Jj c for specimens fro ~ condition A appears significant,

but is probably a reflection of the variaoiity in the flow stress discussed

above. To verify this, representative fractographs were taken from the

surface of the broken C-Ts. Figure 5 shows a picture that is representative

of all of the specimens. The fracture surfaces were similar to other

austenitic stainless steels broken at 4 K No stretch zone was observed

1 296

84
Fig. 5. Typical fracture surface appearance representative of both annealed

and hot rolled C-T specimens: (a) low and (b) high magnification.

between the fatigue precracked region and the J-test region [9]. The J-test

region is covered with dimples. The smaller dimples have nucleated at Nb

and V particles that are uniformly dispersed in the austenite. The long

dimples have nucleated from patches of sigma phase that elongated as a result of hot rolling. The SEM observations in this study support the

previous observations [10,11] that fracture surface features are not

directly related to variations in Jj c.

Simple design considerations can be used to quantify the effect of hot

rolling. The design stress is usually limited by the lower of two factors:

(1) 2/3 yield stress or (2) 1/2 ultimate stress [13]. Looking at the

tensile data in Table 2 for the transverse orientation, the first criterion

is always higher or equal to the second. Criterion #2 would set the design

stress at about 900 MPa, regardless of which condition is chosen.' Uith

similar design stresses, the best choice for the steel's processing is the one with the higher fracture toughness. The data in Table 2 show that

condition D has the same lower-bound value for fracture toughness , about 110

MPaVm, as condition A. There does not appear to be a good reason for

annealing the steel if sufficient control of the hot rolling process is

available. In practice, many steel makers are moving in this direction,

directly quenching austenitic stainless steels after the final pass in the

rolling mill.

CONCLUSIONS

1. Decreasing the hoc rolling temperature increases the 0.2% FS , but

decreases fracture toughness. The calculated QI decreases slightly as the rolling temperature goes down.

2. An annealing treatment after hoc rolling is not necessary to obtain the

best combination of design stress and toughness in this microalloyed

austenicic steel place.

3. The addition of Nb and V to the Fe22Cr- 13Ni- 5Mn-0 . 3N composition

produced carbonitride precipitates that play two important roles in

determining the mechanical properties: first, they helped to maintain a a fine grain size during processing, and second, the precipitates appear

to increase the flow stress and the uniformity of mechanical properties

in this steel, particularly in the annealed condition.

1297

85
4.

The trend llm for fracture toughness as a function of ylaid strength of

annaaiad auscanicic stainless scaals at 4 K doas not axcrapolata to

higher strength levels for hot rolled steals.

ACKNOWLEDGMENTS

This work was partially supported by the Office of Fusion Energy.

Department of Energy. The authors wish to thank M. Riendeau from Rockwell

International and 0. Vigliotti from National Institute of Standards and

Technology for their help in processing the steel and T. Kosa from Carpenter

Technology Corp. for helpful discussions.

REFERENCES

1. R.L. Tobler, D.T. Read and R.P. Reed: in Fracture Mechanics. Thirteenth

Conference, ASTM STP 743 . R. Roberts ed. , American Society for

Testing and Materials, Philadelphia, 1981, p.350.

2. R. Ogawa and J.U. Morris, Jr.: in Advances in Cryogenic EngineeringMaterials , voi 30 . eds. A.F. Clark and R.P. Reed, Plenum Press,

New York, 1984, p. 177.

3. P.T. Purtscher, M. Austin, C. McCowan, R.P. Walsh, R.P. Reed, and J. Dunning: this volume

4. C . A . Perkins: M. Sc. Thesis, Colorado School of Mines, Golden CO, 1987.

5. P.K. Liaw and J.D. Landes: Metal. Trans., vol. 17A, 1986, p. 473.

6. 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.

7. R.P. Reed, P.T. Purtscher, and K.A. Yushenko: in Advances in Cryogenic

Engineer ing- Materia Is , vol. 32 . eds. R.P. Reed and A.F. Clark,

Plenum Press, New York, 1986, p. 43.

8. M.O. Spiedel : in Proceedings of High Nitrogen Steels. HNS--88 . eds. J. Foct and A Hendry. Institute of Metals, London, 1989, p. 92.

9 P.T. Purtscher: JTEVa, vol. 15, Sept. 1987, p. 296.

10 P.T. Purtscher, D.T. Read, and R.P Reed: in Fracture Mechanics:

Perspectives and Directions (Twentieth Symposium). ASTM STP 1020.

eds., R.P. Wei and R.P. Gangloff, American Society for Testing and

Materials, Philadelphia, 1989, p. 433.

11. P.T. Purtscher. R.P. Reed, and D.K. Matlock: accepted for publication

in Proceedings of Materials Research Society meeting on Advanced

Materials, Fracture Mechanics Symposium, Tokyo Japan, June, 1988.

12. N.J. Simon 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. 165

13 S. Shimaraoto, H. Nakajima, K.Yoshida, and E. Tada: in Advances In

Crvoeenic Engineering- Materials . Vol. 32. eds. R.P. Reed and A.F.

Clark, Plenum Press, New York, 1986, p.23.

1 298

86
AUSTENITIC STEELS WITH 9% Cr FOR STRUCTURES AT CRYOGENIC TEMPERATURES

P.T. Purtscher*. M. Austin, C. McCowan, R.P. Reed.

R. P. Walsh, and J. Dunning**

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, CO USA 80303

Advanced Steel Processing and Products Research Center

Colorado School of Mines

Golden, CO USA 80401

**U.S. Bureau of Mines

Albany, OR USA 97321-2198

ABSTRACT

Tensile and fracture toughness tests at 76 and 4 K are performed on austenitic steels containing 9% Cr. The results are comparable to similar

data for AlSI 304 and 316- type steels. Weight- loss experiments show that

these steels will provide good resistance to rust formation. All of the

data indicate that 9% Cr austenitic steels with appropriate alloy additions

would be suitable for structural applications at cryogenic temperatures.

INTRODUCTION

Commercial austenitic (fee) stainless steels (typically AlSI 304 and

316 with 10-12% Ni and 16-20% Cr) are the most common materials for

structural applications at low temperatures. These steels provide good

mechanical properties at the service temperature, a nonmagnetic structure,

and excellent corrosion resistance. The mechanical properties and

austenitic structure are important for devices such as superconducting

magnet cases, but the excellent corrosion resistance afforded by 304 and 316

is not necessary.

Steels with only 9% Cr plus small additions of Ni, Mo, and possibly Cu

and V additions have an austenitic microstructure and corrosion resistance

comparable to 304 in less severe environments. The lower Cr steel did not

present any special problems in melting and hot rolling the material. The

room- temperature mechanical properties are comparable to that of 304

stainless steel [1,2]. No information exists regarding the properties of 9%

Cr austenitic steels at cryogenic temperatures.

Contribution of NIST, not subject to copyright.

4(1\onces m Crvotemc Entmetrmt < Materials). Vol 36

Edited bv R P Reed and F R Fickeu Plenum Press. New York. IWO

1265
Table 1 Chemical Composition of Low Cr Steels (vt.%)

Alloy Cr Ni Mn Si C S 0 P N

X 9.7 8.6 9.5 0.6 2.2 0.047 0.014 0.019 0.006 0.130

Y 8.8 13.7 4.2 0.5 2.0 0.022 0.007 0.025 0.004 0.082

Z 8.7 4.1 14.1 0.6 2.0 0.024 0.007 0.026 0.004 0.100

A 8.5 8.8 9.0 0.4 4.0 0.036 0.010 0.024 0.042 0.077

The goal of our research is to measure the cryogenic mechanical properties and corrosion resistance of several steels with 9% Cr. If the properties are coraparabLe to the commercial austenitic stainless steels, then

9% Cr austenitic steels could be considered for low- temperature service.

Several alloy compositions are studied where the Ni. Mn, and Mo contents are varied while the Cr content is held constant at 9%. Uniaxial tension and

fracture toughness tests in liquid nitrogen (76 K) and liquid helium (4 K) are used to characterize the mechanical properties. Weight- loss experiments

are also run to verify that these compositions would provide adequate

corrosion resistance.

MATERIALS

Four alloys (36 kg heats) were prepared in an induction furnace with

electrolytic -grade alloying elements. The heats were poured in cast iron

molds, 10x10x27.5 cm 3. The ingots were homogenized for 24 h at 1100*C,

upset forged approximately 75%. at 1000°C, and then warm-worked at 700*C into

25 mm thick plates. The plates were given a recrystallization anneal at

1000®C for 45 min and water quenched. The chemical compositions of the four

heats are shown in Table 1.

Table 2. Summary of Weight-Loss Corrosion Tests

Alloy Solution Corrosion Rate

(mpy)

A Glacial acetic acid 0

X Glacial acetic acid < 0.1

Y Glacial acetic acid < 0.1

Z Glacial acetic acid < 0.1

304 S.S. Glacial acetic acid < 20.0

A Phosphoric acid 1.05

X Phosphoric acid 2.15

Y Phosphoric acid 524.0

Z Phosphoric acid 1.4

304 S.S. Phosphoric acid < 20.0

A Distilled water < 0.01

X Distilled water < 0.01

Y Distilled water < 0.01

Z Distilled water < 0.01

304 S.S. Distilled water < 20.0

1 266

88
PROCEDURES

Corrosion testing : Ueighc-loss immersion cases ware conduccad on all

alloys. Coupons 51 nun x 25 mra x 1.5 ram were- cue from annealed place and

polished eo a 120-gric finish. Tescs were conduccad in chree differenc

environmencs : glacial acecic acid ac 80*C, 86% H3PO4 ac 80*C, and discilled wacar ac room cemperacure. The resulcs can be characterized according co NACE classif icacion: fully sacisfaccory resiscance is < 0.5 mm/y (20 rnpy. raillinches/year) ; useful resiscance is less chan 0.5 co 0.5 ram/y; doubeful

ucilicy is 0.5 co less chan 2.5 mm/y; severe attack is > 2.5 ram/y.

Mechanical testing : Uniaxial cension cescs were performed on specimens

oriented transverse co the roLling direction of che plate. The specimens

had a 6.4 mm diameter and a 38 ram long gage length. The flow stress ac 0.2%

plastic strain is defined as the yield strength. The fracture scress is

defined as che true scress at failure.

Fracture toughness tests were conducted on 25 mm thick compact specimens in che T-L orientation according to ASTM standard E 813-87. Tescs were

done in liquid helium (4 K) and in liquid nitrogen (76 K) . The specimens

were fatigue precracked at 76 K and then side-grooved 10% of che specimen

thickness on each side before testing. The details of che exacc cesc techniques are found in an earlier paper [3].

RESULTS AND DISCUSSION

Table 2 shows the corrosion test resulcs. Alloys A, X, and Z were acceptable in all environmencs evaluated. Alloy Y corroded significantly

only in the phosphoric acid solucion. All four compositions would provide

excellent resistance to rusting during che thermal cycling chat cryogenic

structures experience.

Table 3 is a summary of che mechanical test results. In general, the

tensile properties of alloys X. Y. and A were typical of what would be

expected for annealed 316-type steel at cryogenic temperatures. Alloy Z

exhibited significantly lower tensile properties, particularly the elongation and fracture scress values at 4 K . Even at che higher test temperature, the fracture stress is still considerably lower than any of the ocher

chree alloys.

Material characterization : Metallography and X-ray diffraction showed that

che alloys were austenitic after annealing. The microsc rue Cure of alloy Z was heavily faulted, see Fig. 1. but no hep (nonmagnetic) or bcc (magnetic)

martensite was observed before testing. Fractography was performed co document che fracture and deformation behavior of the 9% Cr steels. X-ray

diffraction identified che second phases chat were present after deformation. The x-ray beam contained Cu-Ka radiation with Ka2 stripped away. A

Ge crystal detector measured che diffracted intensities. The relative

volume fractions of each phase were determined by integrated areas under che

peaks found by fitting che data to a Lorenczian profile.

Yield strength is usually che most important mechanical property. A

linear relationship between che N content and yield strength indicates solid

solucion strengthening of che austenite by che interstitial element. Figure

2 shows che yield strengths of the four steels plotted as a function of N

concent and test temperature. The resulcs for alloy Z do not fit the linear

trend drawn through che results from che ocher chree steels. For alloy Z,

che yield strength is apparently controlled by transformation of the aus- tenitic structure co martensite. Excluding tne data for alloy Z, Che slope

of che 4-K data for 9% Cr steels is about 3400 MPa/wc.% N. Over the limited

1267

89
Table 3. Summary of Mechanical Tasc Results from 9% Cr Steels

at Cryogenic Temperatures

Alloy Test

Tenp.

(K)

Yield

Strength

(MPa)

Ultimate % Elong.

,

% R.A.

(Aliform

Strain

(%>

Fracture

Stress

(MPa)

Fracture Toughness

(MPa^m)

Quality

Index x 10* (MPa-MPaJm)

Strength

(MPa)

X 4 851 1510 48.55 3.6 3330 175 148.9

X u 869 1524 45,55 3.8 3380 183 159.0

X 76 668 1302 63.70 45 3284 287 191.7

X 76 672 1334 57,70 45 3518 290 194.9

Y 4 700 1484 48,60 3.1 3690 246 172.2

Y u 720 1491 44.57 4.2 35C& — 177.1

Y 76 577 1259 55,74 42 3344 394 227.3

Y 76 578 1290 57,70 43 3444 425 245.7

Z a 587 1210 21,23 2.4 1578 165 96.9

Z 4 608 1161 14,16 2.4 1389 167 101.5

Z 76 436 1310 52.57 41 2694 317 138.2

Z 76 445 1310 52,57 41 2694 333 148.2

A 4 704 1552 45,57 1.9 3608 227 159.8

A 4 658 1542 43,54 2.8 3218 176 115.8

A 76 528 1336 56,69 45 3236 406 214.4

A 76 544 1364 57,69 44 3546 425 231.2

alloy s 4 0 Mm

Figure 1. Annealed microstructure of alloy Z prior to testing was austenitic with large areas of faults indicated by arrows.

1 268

90
NITROGEN CONTENT, ppm

Figure 2. Tensile yield strength of 9% Cr steels as a function of nitrogen

content and temperature compared to data from ref. 4.

range of N content tested here, the solid solution strengthening of the

lower Cr steels appears to be more effective than for 316-type steels at 4 K

[4] (trend shown on Fig. 2). The yield strength increases at a lower rate,

about 2370 MPa/wt . % N, but with a higher intercept on the y axis of the

plot. Higher yield strengths should be possible in 9% Cr steels if the

nitrogen content is increased, but the race of strengthening may change.

Modified alloy contents to increase the N solubility in austenite or special

melting practices to create a super-saturated austenite with respect to N

[5] may be required co increase the yield strength.

The toughness of alloy Y at 4 K was higher than any of the ocher 9% Cr

steels and alloy Z had the lowest toughness value. Previous research has

shown that the Ni content and yield strength are the most important factors

that determine the fracture toughness of austenitic stainless steels at 4 K. The simplest way to characterize the overall mechanical performance at 4 K

is co calculate a quality index (QI), the product of toughness and yield

strength [6], Spiedel [7] has used the same product co characterize the

performance of steels at room temperature. A higher QI value indicates a better combination of strength and toughness. Previous research has shown

chat increasing the N or Mo content of austenitic stainless steels did not

change the QI whereas increasing Ni content did increase QI [6,8]. In Fig.

3, the calculated QI of the 9% Cr steels is compared to a trend of QI vs Ni

concent for ocher austenitic steels tested at 4 K. The QI values for each

of the 9% Cr steels at 4 K fit the same trend determined for 304LN and

316LN-cype steels. At 76 K, the results for the 9% Cr steels are above the

trend line for 4 K data, but they still show the effect of Ni on the QI

.

Limited data for 304-cype steels at 76 K are also above the*QI-vs-Ni trend

for 4 -K tests [9]. In general, it appears that test results at 76 K give

higher QI values chan test results at 4 K.

Plastic strain at low temperatures transformed part of the austenitic

matrix to hep and bcc martensite in all of the steels. X-ray diffraction

was performed on samples taken from broken, 4-K tensile specimens, approximately 10 mm away from the fracture surface. The volume fractions of each

phase in these steels are shown in Table 4 along with the plastic strain at

thac point. The high hep concent in alloy Z is the most significant point,

1269

91
Figure 3. Calculated quality index vs. Ni content for 9% Cr steels compared

to data from 304 and 316- type steels tested at 4 K.

apparently reducing the properties of the steel when plastically deformed.

Commercial austenitic stainless steels partially transform to martensite,

mainly bcc-tvpe, during cryogenic deformation without any degradation of the

tensile or fracture properties. Tomota et al. [10] have also shown that a large volume fraction of hep martensite can reduce ductility in tensile

tests

.

Fractography on the broken compact tension specimens to documented the

micromechanisms of fracture which occur in the 9% Cr steels. No stretch

zones were observed between the fatigue precracked region and the J-test

region, similar to the appearance of commercial austenitic stainless steels

tested at cryogenic temperatures [11]. Figure 4 shows an example of the

dimpled rupture morphology that was found on the J-test region of specimens

from alloys X and Y tested at 4 K. Similar features were observed for

steels X, Y. and A when tested at 76 K. Voids form at non-metallic inclusions in the austenite to start the fracture process.

Table 4. Phases Present after deformation of 9% Cr Steels at 4 K

Alloy

True - Strain fee bcc hep

X 0.39 40 34 26

Y 0.48 29 71 0

Z 0.14 27 8 64

A 0.33 36 40 24

1270

92
Fig.4.SEM micrograph of Che fracture

surface of a compacc specimen

from alloy X tested at 4 K.

Fig.S.SEM micrograph of the fracture

surface of a compacc specimen

from alloy Z tested at 4 K.

For alloy Z, the fracture morphology appears crystallographic; see Fig5. Cross sections through the fracture plane were examined and compared to

the appearance of the fracture. Apparently, the fracture starts at the hep

martensite places chat are present ahead of the crack tip rather chan at

nonmecallic inclusions. The fracture morphology for alloy A tested at 4 K was also crystallographic, but with some evidence of void nucleacion at

nonmecallic inclusions.

It has been suggested that martensite formation during mechanical

testing is responsible for transformation* induced plasticity (TRIP) where

the mechanical properties are greatly enhanced at certain test temperatures

[12]. No significant improvement in any of the mechanical properties of

these 9% Cr steels has been observed at 4 or 76 K. On the contrary, the

poor mechanical properties of alloy Z at cryogenic temperatures are directly

related to the presence of hep martensite. If there is a TRIP effect in our alloys, it must lie at some temperature above 76 K.

CONCLUSIONS

1. With appropriate alloying, 9% Cr steels have mechanical properties which

are equivalent to commercial stainless steels at cryogenic temperatures. The

Ni content is a key ingredient for low- temperature toughness. The N content

is most important for controlling the yield strength.

2. All four of the 9% Cr steels tested here have sufficient corrosion

resistance for currently anticipated cryogenic structural applications.

3. The formation of hep martensite in 9% Cr austenitic steels can reduce

tensile and fracture properties at cryogenic temperatures.

ACKNOWLEDGMENTS

The help of Luming Ma, a guest work at NIST on leave from the Chinese

Institute of Metals, Shenyang, China in performing many of the tensile tests

for this program is gratefully acknowledged. The research was partially

supported by DOE Office of Fusion Energy.

1271

93
references

1. S. Floreen, Metal. Tran*., 13A:2003 (1982).

2. J.S. Dunning, J.M. Oh, and J.C. Rawers, in: Alternate Alloying for

Environmental Resi stance, edited by C.R. Smollk and S.K. Banerji,

Metallurgical Society of AIME, Uarrendale, PA (1987), p. 3.

3. R.L. Tobler, D.T. Read, end R.P. Reed, in: Fracture Mechanics . Thirteenth Conference . ASTM STP 743, R. Roberts, ed. , American

Society for Testing and Materials, Philadelphia (1981) p. 350.

4. N.J. Simon 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. 165.

5. R.P. Reed, Journal of Metals, 41:16 (1989).

6. R.P. Reed, P.T. Purtscher, and K.A. Yushenko, in: Advances in Cryogenic

Engineering. Materials. Vol. 32 . eds. R.P. Reed and A.F. Clark,

Plenum Press, New York (1986), p. 43.

7. M.O. Spiedel , in: High Nitrogen Steels, HNS -88 . eds. J. Foct and A.

Hendry, Institute of Metals, London (1989), p 92.

8. 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.

9. P.T. Purtscher and R.P. Reed, in: High Nitrogen Steels--HNS 88 . eds. J.

Foct and A. Hendry. Institute of Metals, London (1989), p. 189.

10. Y. Tomoto, M. Strum, and J.U. Morris, Jr., Metal. Trans., 17A:537

(1986) . 11. P.T. Purtscher. JTEVA, 15:296 (1987).

12. S.D. Antolovich and B. Singh, Metal. Trans., 2:2135 (1971).

1272

94
MICROMECHANISMS OF DUCTILE FRACTURE AND FRACTURE

TOUGHNESS IN A HIGH-STRENGTH AUSTENITIC STEEL

P.T. Purtscher

Materials Reliability Division

National Institute of Standards and Technology

Boulder, CO USA

The micromechanisms of ductile fracture in a high

strength austenitic stainless steel have been studied as a

function of test temperature and stress state. Two different micromechanisms are important in this steel: void

formation that leads to microvoid coalescence (MVC) and

shear fracture. The differences between the two micromechanisms is related to when strain localization occurs.

For MVC, strain becomes localized between voids that nucleate when a critical stress is reached during uniform

deformation. For shear fracture, strain becomes localized

across the deforming section before any significant volume

fraction of voids have nucleated, that is, before the stress

for nucleation is reached. In uniaxial and biaxial stress

states, shear fracture becomes the important micromechanism

of failure as the test temperature decreases because the

stress for void nucleation increases for lower test temperatures. In the triaxial stress state associated with fracture toughness testing, MVC is the dominant micromechanism

observed on the fracture surface of specimens tested between

295 and 4 K. Observations from interrupted fracture toughness tests show that void nucleation from large, scattered

particles has little influence on the measured fracture

toughness. Void nucleation, growth and coalescence at

small, closely spaced microalloy particles occurs at an

applied value of J less than J,

c

, the critical value that is

associated with the measured fracture toughness of the

steel. Observations and preliminary calculations suggest

that fracture toughness is related to void nucleation ahead

of the crack where the hydrostatic stress associated with

crack tip blunting is a maximum.

95
1 . 0 INTRODUCTION

The effect of temperature on the fracture of iron

alloys depends on the crystal structure of the dominant

phase. In ferritic steels, body-centered-cubic struct re,

the normal fracture behavior at room temperature (295 K) is

ductile. As the test temperature is decreased, a brittle,

cleavage fracture occurs at a critical temperature, commonly

called the ductile-to-brittle transition temperature ( DBTT )

.

Fracture of ferritic steels at temperatures near the DBTT

has been studied extensively [1] and is not considered here.

In contrast, most austenitic stainless steels, face-centered-cubic structure, do not undergo any transition to

cleavage fracture as the temperature is decreased to temperatures as low as 4 K [2]. Only ductile fracture is important for ductility and toughness in austenitic stainless

steels.

Austenitic stainless steels are commonly used for

structural applications in severe environments, for example,

superconducting magnet cases, liquefied gas storage vessels,

and rotors for turbine generators at elevated temperatures.

All of these applications require good mechanical properties

such as strength, ductility, and toughness, in addition to

excellent corrosion resistance.

Recent alloy development programs have improved the

mechanical properties of austenitic stainless steels [3-4].

The strength can be raised by additions of nitrogen and

careful control of grain size during processing. The basic

mechanisms that control tensile ductility are not veil known

for austenitic stainless steels. The relationship between

ductility and toughness is vague because there are different

ways to evaluate toughness. In this study, toughness is

discussed in terms of fracture toughness. Fracture toughness is a material property that describes the conditions

under which an initially sharp crack begins to grow. The

basic framework for defining fracture toughness was proposed

in the late 1950's by Irwin [5]. Fracture toughness is more

valuable than other measures of toughness because it directly relates to the performance of a material in service.

Fracture toughness is the only type of toughness that is

considered in this work.

The micromechanisms related to toughness and ductile

fracture in an austenitic stainless steel are the subject of

96
this work. Much has been written on this subject, and

several good review articles and books have recently been

published [6-10]. These references provide a thorough

description of the present state of knowledge regarding ductile fracture, and only background information pertinent to

the micromechanisms of ductile fracture in steels are discussed here. The state of the art in toughness testing and

modeling is reviewed briefly with respect to how the micromechanisms of ductile fracture relate to toughness. The

next section summarizes preliminary work on toughness of

austenitic stainless steels at cryogenic temperatures. At

the end of the introduction, the hypothesis for this report

is presented along with a proposed experimental plan.

1.1 Ductile fracture

There are many definitions of ductile fracture. All

agree that ductile fracture requires plastic deformation.

For example, Dieter [11] defines ductile fracture as

"fracture occurring with appreciable gross plastic

deformation." In this work, materials that fail after

localized plastic deformation are called "ductile" regardless of the micromechanisms that lead to fracture. A micromechanism is a physical process that operates as a result of

plastic deformation. The micromechanisms that cause ductile

fracture are those that produce localized deformation.

Ductile fracture under tensile loading can be

classified according to three fundamental micromechanisms

[11] shown in Fig. Is (a) necking to a point, (b) shear,

and (c) pore formation or void nucleation (referred to as

microvoid coalescence or MVC) . Neck formation and fracture

at a point (Fig. la) occur when plastic deformation is

localized by neck formation and continues until the macroscopic cross sectional area reduces to a point. Void nucleation is not required for fracture by necking to a point.

This first type of ductile fracture is typical of soft, pure

materials and/or, materials in high-temperature tests.

Shear fracture (Fig. lb) occurs when plastic strain is

localized into a shear band oriented at 45° to the tensile

axis and is characteristic of single crystals in tension.

Again, void nucleation is not required for shear fracture.

Void nucleation, growth, and coalescence are often observed

on the fracture surface of shear fractures, but the MVC

features form after strain is localized in the shear band.

97
Figure 1. Schematic diagram of micromechanisms of ductile fracture: (a) necking, (b) shear, and

(c) void formation.

MVC (Fig. lc) is a three step process. First, voids

nucleate inside the material. After nucleation, void growth

starts and continues to a critical point where voids coalesce to form a fracture surface. In steels, nucleation

normally occurs at second phase particles in the matrix when

a critical stress is reached. The voids can form either by

decohesion of the interface between the particle and the

matrix or by particle cracking. Void growth is related to

plastic deformation in the intervoid regions. The ductile

behavior is commonly observed in both ferritic steels,

strained at a temperature above the DBTT, and austenitic

steels.

To summarize, fracture by necking to a point is the

simplest case. No voids are required to form and no shear

98
bands form. Shear fracture occurs when strain is localized

into a shear band before void nucleation. MVC occurs when

void nucleation occurs before strain is localized into a

shear band. After primary void nucleation, fracture of the

ligaments between voids is caused by one of three mechanisms: a neck forms and deforms to a point, strain is localized into a shear band, or a second set of smaller voids

form.

Because many engineering materials fail by a ductile,

MVC mode in service, MVC has been studied extensively [6-

10,12-15]. The most common method used to study MVC has

been to measure the area fraction of voids found on the

polished cross section of a deformed tensile specimen as a

o 4

O

>

o 3

<

LU

tr

<

iS

o

2

0

* * * 1 0.6 0.5 0.4 0.3 0.2 0.1 0.0

TRUE THICKNESS STRAIN

Figure 2. Plot of area fraction of voids vs. true

strain for sheet tensile specimens of a dualphase steel tested at room temperature [ 12 ]

.

99
function of strain. Figure 2 shows the results of such a

study of necking and ductile fracture in a dual-phase steel

[12]. The first voids nucleate at a relatively small

strain. The area fraction of voids increased only after the

specimen was strained past the uniform strain. The void

nucleation strain can be defined as the strain where voids

were first observed. In this case, the nucleation strain is

a small fraction of the fracture strain. The absolute value

of nucleation strain will be a strong function of the resolution of the technique used to observe nucleation. Conversely, the strain for void nucleation can be operationally

defined by extrapolating the void area fraction to 0, as

shown in Fig. 2. In this case the void nucleation strain is

a large fraction of the fracture strain.

The increase in void area with strain above the nucleation strain is due to nucleation of new voids as well as

growth of existing voids. Individual voids start to coalesce to form macrovoids, which then can be described as a

crack. Additional strain may be required to fracture the

specimen after void coalescence as observed by Park and

Thompson [13]. Growth and coalescence are concentrated in

the necked region of the specimen and occur over a small

increment of remote, applied displacement in the test.

Models for void nucleation are normally based on void

decohesion of the interface between the matrix and secondphase particles. Decohesion occurs when a critical strain

or stress is reached at the interface. Argon and Im [14]

applied a critical stress criterion to void nucleation;

theirs is the most generally accepted approach. They

suggest that decohesion occurs when the stress at the

interface, a ]f reaches a critical value, o

c

. The value of o. is defined as

o i + 0T' ( 1 ) where * the equivalent flow stress of the material and o T = the imposed hydrostatic stress.

Beremin [15] applied the Argon model for void nucleation to the formation of voids at MnS stringers in A508

Class 3 steel, a C-Mn steel typically used in construction.

Argon* s nucleation criterion did not yield a consistent

value of o for different test temperatures and test orientations. The following criterion for o

c fitted all the

experimental observations of void nucleation:

100
o (2) c = o, + 0(0^ - o 8 ) , where o 1 is the maximum principal stress and o

Q

is the yield

stress. The constant C is a fitting parameter that was

equal to 1.6 for longitudinal and 0.6 for short transverse

orientation.

Modeling of void growth and coalescence [16-18] usually

assumes that nucleation and growth are separate events and

that void growth is proportional to plastic strain. The

extent of growth depends on the stress state and material

properties. As Fig. 2 shows, voids start to grow in the

necked region and the specimen fractures when the void area

fraction approaches 0.1. Thomason [19,20] has suggested a

different approach to model void growth that considers the

effect of void nucleation on growth as well as the effects

of strain localization in the matrix, which depends on the

stress state.

1.2 Fracture toughness

1.2.1 Description and measurement of toughness

There are two basic methods to describe toughness, the

local crack opening displacement (6)

,

or the remote, applied

energy required to open the crack [21]. The applied energy

methods can be described by a stress intensity factor (K)

when the plastic zone size is much smaller than the dimensions of the specimen (linear-elastic, small-scale yielding) , and by the J- integral (J) when the plastic zone is

comparable to the specimen dimensions (elastic-plastic, or

large-scale yielding)

.

Methods to measure toughness have been developed to the

point where standard procedures are well accepted and the

external factors, like specimen design and loading rate,

that affect toughness are well established. In the United

States, the energy methods are often preferred over 6

testing. The most common measurement methods are standard

ASTM tests, E-813 for large-scale-yielding and E 399 for

small-scale-yielding. Both methods use small, fatigue

precracked specimens with a gage mounted at a convenient

location to monitor the specimen displacement while the load

is applied. A load vs. displacement curve (P-u) is recorded

to characterize the test. The transition from small-scale

yielding to large-scale yielding can occur for a given steel

101
by decreasing the specimen dimensions or increasing the test

temperature.

The critical point for toughness measurement in both

large- and small-scale yielding is represented by a

significant change in the compliance or stiffness of the

specimen. In the case of large-scale yielding, the change

is more gradual and the crack grows in a stable manner. For

the E-813 procedure, periodic unloadings are performed to

monitor the change in compliance. The criterion that defines toughness is a specified decrease in compliance. In

small-scale yielding, the compliance drops rapidly and the

critical point can be determined by inspection of the loaddisplacement (P-u) curve. Therefore, a key question is,

"How do the micromechanisms of ductile fracture change the

compliance of the test specimen?"

A generally accepted correlation between micromechanisms and compliance for a "ductile" material taken from

the literature [22] is presented in Fig. 3. The P-u curve,

Fig. 3a, is linear over a short range of displacement (elastic behavior) , and the curve becomes nonlinear (plastic

behavior) when the material yields. The test is stopped

briefly at regular intervals and unloaded to check the

compliance of the specimen. At point 1 in Fig. 3, the

elastic stresses at the crack tip are higher than the flow

strength of the bulk material. To relieve the stresses near

the tip, the crack blunts (changes shape at the tip) , and a

plastic zone forms around the tip. The shape change produces a small increment of crack growth, called a stretch

zone (SZ)

,

from the original fatigue crack. Voids will

nucleate and grow in the material as soon as the plastic

zone engulfs a nucleation site, Fig. 3b. A small change in

compliance is measured because the crack has grown slightly.

A data point is then plotted on the energy vs. change in

crack length (J-*a) curve. Fig. 3c.

After unloading at point 1, the specimen is reloaded

and more nonlinear behavior is seen in Fig. 3a and more

unloadings are made. As long as the nonlinear behavior is

due to plasticity, the crack tip in Fig. 3b will continue to

blunt and the size of the plastic zone will increase. Any

voids close to the tip can grow while new void-nucleation

sites are activated in the plastic zone. The SZ continues

to grow relatively slowly and requires considerable energy

per unit crack growth. Each unloading produces a data point

in Fig. 3c. At applied J < Jlc , the points lie along a

102
LOAD-LINE DISPLACEMENT

(b)

(c)

Figure 3.

>*

CD

cc

LDZ

LU

'Ic

Blunting

..-Ar'*"

/3

4

X2

0 *

*-7- *

Tearing

Delta a

Schematic diagram showing the correlation

between micromechanisms of ductile fracture

and fracture toughness: (a) load-displacement curve, (b) crack tip profile, and (c)

energy vs. crack growth curve [22].

103
line, referred to as a blunting line with a slope = dJ/da.

From experience, dJ/da can be approximated by the properties

of the material in a uniaxial tension test.

At some critical point, labeled 4 in Fig. 3 for this

example, the growing voids ahead of the blunted tip impinge

on each other, so there is a continuous surface connected to

the fatigue crack and SZ in Fig. 3b. This change in crack

length produces a significant change in compliance. Point 4

in Fig. 3c is below the extension of the blunting line. At

each successive unloading along the P-u curve, more of the

growing voids link up with the fatigue crack. The crack

growth rate at this point in the test is called the tearing

line or regression line in E-813 and can be characterized by

the tearing modulus (T)

.

The intersection of the blunting

and tearing lines defines J

lc , the parameter that characterizes the toughness of the material. The local displacement values, 6 and SZ, at this critical point could also be

used as characterizing parameters.

If the same ductile material is tested again at the

same temperature, but with a sufficiently large specimen,

small-scale yielding could exist. The small-scale yielding

may promote a different failure mechanism than the observed

mechanism under large-scale yielding. If no transition infailure mechanism is observed, the correlation between

external loading and micromechanisms should be similar to

the events described in Fig. 3, only the processes occur

inside a smaller plastic zone.

Small-scale yielding is often called "brittle” or

stress-controlled fracture because the specimen can break

after little or no measurable yielding. Large-scale yielding is called "ductile" or strain-controlled fracture

because there is obvious deformation before the specimen

breaks. In both large- and small-scale yielding, the fracture surface can have the same microscopic features (MVC)

.

The macroscopic roughness of the fracture surface and the

shape of the P-u curves do not correlate directly to the

microscopic fracture mode. The terms "ductile" and "brittle" are misleading in this context because they are relative terms that describe the scale of the plastic zone

compared to the specimen dimensions, not the micromechanism

of fracture. For the two cases, small-scale yielding and

large-scale yielding, the material behavior is different,

but both can have the same measured toughness. Jlc

and

are related by the elastic properties of the material. This

104
equality is the major reason to use fracture mechanics.

Small specimens can be tested to predict the behavior of

much larger structures, provided that the failure mechanism

is the same in both the large structure and small test

specimen. A more thorough discussion of this subject is

given by Banerjee [23].

1.2.2 Modeling of toughness

The stress and strain distributions ahead of the crack

tip form the basis for modeling toughness. Many researchers

have modeled the stresses and strains around the crack tip

of a fracture mechanics specimen as a function of 6 with

analytical and numerical approaches [24-27]. The finite

element method (FEM) results of McMeeking [24], shown in

Fig. 4 for a low-strength, nonhardening material, are the

most often referenced. The radial stress is highest ahead

of the crack with a maximum value of 3 to 4 times the tensile yield strength at a distance of 26. The strain is

highest at the crack tip and decreases exponentially with

the distance from the tip, but is always highest at an

angle, 0, of n/4 with respect to the original crack tip.

McMeeking concluded from his results that if the direction

of initial crack growth was straight ahead of the crack tip,

then the maximum stress controlled the fracture. If the

crack grew initially at a 0 = n/4 to the tip, then the

strain was more important. The numerical results are in

good agreement with analytical results of Hutchinson [25],

and Rice and Rosengren [26] (called the HRR strain and

stress fields) , and the Rice and Johnson [27] slip-line

analysis.

Shoji et al. [28] and Luo et al. [29] measured the

strain distribution around a blunted crack tip in ferritic

steels at room temperature by different techniques. In both

cases, the results agree with the FEM results for the region

directly ahead of the crack tip. However, the strain was

not a maximum at 0 = ir/4 to the crack plane. The highest

strains were found directly ahead of the crack tip. Also,

no voids were observed in the matrix ahead of the crack tip.

The simplest approach to modeling [30] assumes that

toughness is the product of the average flow stress in the

plastic zone (o Y

) , the local strain to failure (e f) , and a

characteristic distance (L) related to the spacing of weakly

bound inclusions where voids can nucleate easily. One form

105
Figure 4 . Plot of FEM results showing the distribution

of stress and strain ahead of the crack tip

in a nonhardening material [24].

of the equation with Jlc as the toughness parameter is,

J

lc = o Y x e f x L. (3)

The main difference between models is the way in which

e, is determined. Some models assume that the fracture

strain measured in a plane strain tensile test is the same

as e f in the model for toughness. Another approach is to

assume that the fracture surface morphology is a measure of

e

f

. Ritchie and Thompson [30] assume the microroughness is

a measure of e f while Garrison [31] suggests that the ratio

of inclusion size to void size is related to e f

. These

types of fracture surface observations assume that the

growth of the largest voids are related to toughness and

that void nucleation does not contribute.

A recent experimental study of fracture mechanisms [32]

measured the mechanical properties of ASTM A533B steel as a

106
function of the size of pre-existing voids in the matrix.

The results showed that the presence of pre-existing voids

of any size reduced the tensile and toughness properties

significantly compared to the properties with no preexisting voids. The fracture surface appearance suggests

that void coalescence starts almost immediately when voids

are present prior to plastic strain. A similar experiment

with silver [33] showed the same results. The strain appeared to concentrate in the ligaments between pre-existing

voids with little uniform strain. The energy required to

nucleate voids was a large fraction of the tensile ductility

and contributed significantly to toughness.

Two different research groups [34,35] have tried to

incorporate Beremin's criterion for void nucleation into a

model for toughness. In both cases, the void-nucleating

particles were large, elongated inclusions in a low-strength

(< 500 MPa) ferritic steel matrix. The nucleation criterion

was easy to achieve in the triaxial stress state at the

crack tip, so resistance to void nucleation did not contribute significantly to toughness.

1.2.3 Toughness correlations:

The toughness of many engineering materials has been

measured according to the standard test methods described

above and correlated to tensile properties. A summary of

toughness data at room temperature for different wrought

engineering alloys is plotted as a function of tensile yield

strength in Fig. 5 [36]. Data for austenitic stainless

steel at 4 K [37] have been added to Fig. 5. All the alloys

show a similar trend in which toughness decreases as yield

strength increases. Scatter within the band for each material is an indication of other factors that influence toughness. The important "other” factors are the melting practice and residual alloy content that determine the volume

fraction, size, distribution, and type of inclusions.

Alloy development programs have improved the properties

of structural steels over the years by optimizing the

factors that influence the mechanical behavior of steel.

Some results are summarized in Fig. 6 [38]. The strength

and toughness have increased since 1960, moving the curve

towards the upper right comer of the plot.

Speidel [38] suggests that the product of yield strength and

toughness, defined here as the quality index (QI) , appears

107
YIELD STRENGTH, MPa

Figure 5. Plot of fracture toughness vs. yield strength

for high-strength structural alloys from [36]

and similar data for austenitic stainless

steels at 4 K [37].

108
Fig. 6. Plot of toughness vs. yield strength for high

strength structural steels showing how the

properties have improved since 1960 [38].

to be a valuable indicator of improved mechanical properties. Speidel found that an austenitic, Cr-Mn-N steel had a

higher QI than any other structural steel at room temperature .The QI can also be used to quantify the effect of

thermal-mechanical processing and changes in alloy composition on mechanical properties of austenitic stainless

109
steels at cryogenic temperatures [39,40]. A summary of the

QI at cryogenic temperatures for austenitic stainless steels

with different compositions and processing is shown in Fig.

7 as a function of Ni content [41]

.

A higher temperature

results in a lower yield strength, but the increase in

toughness produces a larger QI . At a constant temperature,

a higher Ni content has little effect on yield strength, but

a significant increase in toughness.

NICKEL CONTENT, wt%

Figure 7. Plot of the quality index of austenitic stainless steels at 4 and 76 K as a function of

Ni content [41]

.

110
1.3 Preliminary work on toughness of high strength

austenitic stainless steel at low temperatures

The most significant observation regarding austenitic

stainless steels at cryogenic temperatures is that these

steels fracture without an observable SZ on the fracture

surface [42]. The fracture surfaces of two austenitic

stainless steels and one ferritic steel are shown in Fig. 8.

Each has nearly the same measured toughness, yet the surface

morphology at the tip of the fatigue crack in the austenitic

stainless steels show no sign of a SZ. A separate region,

the SZ, is easily seen between the fatigue crack and the MVC

features of the toughness test in the ferritic steel.

This apparent anomaly was explained in a paper in terms

of the micromechanisms of ductile fracture [43]. In order

for a SZ to form, the material directly ahead of the crack

tip had to be strained significantly before voids nucleated.

The distribution of strain around the crack tip of an austenitic steel (AISI 304) at 76 K is shown in Fig. 9. At the

crack tip, the strain is concentrated at 9 = n/4 to the tip.

A crack forms at the crack tip when the applied J is less

than JJc. Separate voids were not observed ahead of the

crack tip.

Another key observation was that the initial crack

growth in austenitic stainless steels at cryogenic temperatures was straight ahead of the fatigue crack even under

large-scale yielding. McMeeking [24] predicted these observations from his FEM analysis when a critical stress criterion controlled fracture. Ferritic steels typically showed

more angled crack growth in the direction of the maximum

strain. For both the ferritic and austenitic stainless

steels, the micromechanisms for ductile fracture were limited to a region approximately 25 around the crack tip. This

region was defined as the process zone (PZ)

.

A simple model to describe the stress intensity factor

required to nucleate voids ahead of a blunt crack was described by Purtscher et al. [44]. This approach assumed

that void nucleation occurred when the interfacial stress

between the matrix and inclusion reached a critical level,

o

c

, that was equal to the maximum applied stress plus the

local stress due to dislocations. For austenitic stainless

steels at 4 K, the applied stress was related to the yield

strength of the steel, and the local stress was controlled

by the strain that was related to the Ni content of the

111
112

Fig.

8

Comparison

of

the

fracture

surface

from

different

steels:

AISI

316LN

tested

at

4

K,

(b)

AISI

308L

tested

at

76

K,

and

(c)

HY

80

ferritic

steel

tested

at

295

K.
3 m oo

014J.C 0 C/1 C T3 £00

CT£ £ 3 0 O £ T3 £00£

<4-1^0 O ft £

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•H ^ £ •• £ « — *0 <0 T3

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C/1

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C/1 *H 0 H (C •H < P £ £ a£ c/i <o o p 00

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cr>

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113

grain
Hot Deformation Sequence

(

1000°C - 1150°C

Rolling

Figure 10. Schematic diagram showing the hot deformation

processing used to produce a plate section

from the as-received bar. The dark crosshatched region in the center of the cross

section represents the region that contained

sigma phase in the starting microstructure.

114
steel. The model was extremely sensitive to the assumed

value of the local and macroscopic stress. Given best

estimates of the stress terms, the model predicted that

nucleation accounts for the majority of the measured tough

ness for austenitic steels at strengths above 600 MPa.

1.4 Purpose and Experimental Program

The purpose of this report is to determine the effect

of temperature and stress state on the micromechanisms that

cause ductile fracture in an austenitic steel. Observations

of the micromechanisms of ductile fracture are correlated to

the macroscopic stress and strain conditions in terms of two

different criteria, Eqs. 1 and 2, as proposed by Argon et

al. [14] and Beremin [15] . Critical conditions for ductile

fracture in simple stress states is related to ductile

fracture in a complex stress state present in the fracture

toughness test. The toughness is modeled with a MVC model

(void nucleation, growth and coalescence) by the the matrix

plastic flow properties and the distribution of second-phase

particles in the matrix.

Other researchers [30-32,34,35] have studied the effect

of matrix plastic flow properties and second-phase particle

distributions on toughness. In the past, the critical event

was interpreted to be the growth of voids from the largest

second-phase particles. Void nucleation was assumed to contribute little to the total energy absorbed.

The first step to understand ductile fracture and

toughness better is to quantify the role of void nucleation.

In the past, very little quantitative data have been

generated. Clearly, nucleation is the first stage in the

overall process and must be accounted for in any generally

applicable theory. It is not clear if the same criterion

required for void nucleation in a tension test can be

applied to different stress states and test temperatures.

Other key questions relate to the role of second-phase

particles. How can the critical inclusions involved in

nucleation be determined? Is the average inclusion spacing

the important parameter, or is the actual distribution

important for determining the characteristic distance?

Clusters of inclusions have been known to promote

cleavage fracture in ferritic steels [1]. The characteristic distance, L, used in most models is always equated with

the spacing of void-nucleating inclusions. In practice

115
[32], the model is assumed to fit the data, and a value of L

is calculated from the measured toughness. L is then equal

to some multiple of the inclusion spacing, sometimes less

than 1 or other times 6 to 7 times the spacing. L is used

more as a fitting parameter with no physical significance.

A different approach to the idea of a characteristic distance for toughness testing could provide insight.

The final point addressed in the research deals with

the material properties of austenitic stainless steels. The

influence of different constituents in the microstructure

are described in terms of the observations.

2 . 0 MATERIAL

The steel chosen for this study is a commercially

available nitrogen-strengthened austenitic stainless steel

with the composition shown in Table 1. The steel, called

22-13-5 in this report, was received in the form of a round

bar, 76 mm in diameter. Perkins [45] studied the hot deformation behavior of the same heat of 22-13-5 steel.

The microstructure of the as-received bar was initially

characterized by Perkins [45]. The center of the bar contained a large fraction of second-phase particles distributed in the austenitic matrix. The elongated, second phase

was identified as sigma phase. The hardness varied from R

A

116
61 near the outer surface to RA 54.5 in the center of the

bar. Five tensile specimens were machined from the bar,

evenly spaced through the diameter and parallel to the

rolling direction, and tested at 4 K to characterize the

starting material. The results of these initial tests are

shown in Table 2. The 4-K yield strength and elongation are

higher around the edge of the bar and decrease in the center

of the bar. The tensile properties and hardness reflect the

Table 2 Tensile Test Results

22-13-5 Bar

at 4 K for

YS UTS Elongation Reduction

Location (MPa} (MPa^ % in Area %

Edge 1540 1877 24 34

Mid-radius 1212 1810 16 30

Center 1107 1794 13 27

Mid-radius 1171 1792 21 37

Edge 1503 1858 28 40

influence of the large volume fraction of sigma phase in the

center of the bar.

Sections of the as-received bar were rolled into

plates, approximately 28-mm thick, 105-mm wide, and 63-cm

long. The hot deformation sequence for processing the asreceived bar is shown in a schematic diagram in Fig. 10.

The round bar was first deformed into an elliptical shape in

a steam hammer and then rolled into a plate with a

rectangular cross section. The lighter cross-hatched region

of the cross section represents regions free of sigma phase.

The darker cross-hatched region in the center of the cross

section represents the region that contained large volume

fraction of sigma in the starting microstructure. The

temperature for final hot rolling was varied between 1150

and 950 C. All the specimens in this research program received a final annealing treatment at 1150 C for 1 to 2 h,

followed by a water quench.

In a separate program, specimen blanks were machined

from the hot-rolled material to evaluate the effect of the

different processing histories on mechanical properties at

4 K. Those results were reported eleswhere [40]. Lower

117
hot-rolling temperatures resulted in higher 4-K yield

strength, but slightly lower tensile ductilities and toughness .

3 . 0 EXPERIMENTAL PROCEDURES

3.1 Material Characterization

Specimens for light microscopy were taken from both

deformed and undeformed regions of the various test

specimens. The samples were prepared by standard procedures

for mounting and mechanical polishing. The samples were

then etched with different etchants to reveal various

microstructural constituents [46]. Sigma phase was

identified by an electrolytic KOH etch that attacks only

sigma and delta ferrite; sigma is differentiated by color

(sigma is stained darker than delta ferrite) . The grain

size and deformation structure were delineated by an electrolytic nitric acid and water etch.

Scanning electron microscopy (SEM) was used for

fracture surface observations and inclusion counts on the

cross sections. All inclusions with average diameters

greater than 0.5 pm were counted by a particle recognition

program (PRC) . Each analysis took between 2 and 3 h to

count individual frames at 500x, and to measure the shape

and size parameters associated with each inclusion.

Transmission electron microscopy (TEM) of thin foils

was used to observe the distribution of sub-micrometer size

particles and to evaluate deformation structures as a

function of deformation temperature. An electrolyte

(containing 5% perchloric acid, 75% ethyl alcohol, and 20%

glycerin) was used at room temperature with a current of

about 80 mA. Four different deformed conditions were observed, two at 295 K (strain = 0.02 and 0.05) and two at 76

K (strain = 0.03 and 0.12).

Microhardness measurements at 50- and 200-gram loads

were used to characterize the annealed and deformed material

on a fine scale. Around the crack tip, microhardness

readings were taken in a cylindrical coordinate system

centered at the tip of the fatigue precrack. The parameter

r is the distance from the center. The angle 0 refers to

the angular displacement from the plane parallel to the

118
fatigue crack with 0=0 directly ahead of the crack and 6 =

rc/2 and -it/2 perpendicular to the crack.

3.2 Mechanical testing

3.2.1 Uniaxial tension (uniaxial stress state)

Blanks, 13 -mm square, were removed from the top and

bottom of the plates in both transverse (T) and longitudinal

(L) orientations with respect to the rolling direction.

Specimens with a 6.25-mm-gage diameter were machined with

either a 45- or a 28-mm-gage length. The specimens were

tested in a screw-driven test machine with a 100-kN load

capacity. The displacement was measured by a special extensometer for cryogenic service [47]. The initial cross-head

rate was 0.05 cm/min for tests at 295 and 76 K. For tests

at 4 K, the initial cross head rate (CHR) was 0.02 cm/min.

For most tests, the CHR was increased to 0.1 cm/min after a

strain of about 0.05. When the fracture stress and hydrostatic stress concentration factor are reported, the CHR was

held constant at 0.05.

The load and displacement data for each tension test

was recorded continuously on an AT-compatible data acquisition system. The data were analyzed to determine the yield

strength (YS) , ultimate strength (UTS) , percent elongation

(%EL) , and reduction in area (RA) • YS is defined as the

flow stress at a plastic strain of 0.002. UTS is defined as

the maximum load divided by the original cross sectional

area. %EL is the change in change in length of the uniform

gage section divided by the original length. RA is the

change in cross-sectional area at the point of fracture

divided by the original area. The strain hardening rate was

obtained by first smoothing the stress and strain data with

a nine-point average, and then calculating a slope at each

point on the curve that represents true stress vs. strain.

The fracture stress for those specimens tested to failure at

a constant CHR is defined as the load at the point of fracture divided by the final cross-sectional area.

Flat tension specimens with a 25-mm-long gage section

were machined transverse to the rolling direction according

to the design shown in Fig. 11a. These specimens were

tested to different strain levels at 295 and 76 K for later

evaluation by TEH; these tests were run in the same test

machine used for uniaxial tension testing.

119
3.2.2 Plane-strain tension testing (biaxial stress state)

Tests were conducted at three test temperatures and in

two ways: to the point of fracture and interrupted at a

given strain. The flat, plane-strain specimens had a gage

section that was 2 x 25 x 0.75 mm, see Fig. lib. The design

was used by Corrigan [48] to produce a biaxial stress state

and maximum constraint in the gage section. The specimens

were machined with a transverse orientation with respect to

the rolling direction so that the width of the specimen, the

direction of zero strain in the specimen, traversed the

thickness of the plate.

The specimens were tested in a servo-hydraulic controlled test machine operated in stroke control. The displacement rate was 0.02 cm/min, which corresponded to an

elastic loading rate of 1.44 kN/min. The load and change in

thickness of the specimen were monitored during the test

with an X-Y recorder. The through-thickness strain, e t , was

calculated according to the equation,

€ t = In t</ ti, (4)

where to = initial thickness and t,* = instantaneous

thickness

.

3.2.3 Fracture toughness testing (triaxial stress state)

Fracture toughness tests were conducted according to

ASTM standard procedure E 813-81 at the three test temperatures. The unloading compliance method was used to monitor

crack growth. The 1981 standard analysis calculates Ji c from

the intersection of the blunting line, with a slope that is

twice the average of the YS and UTS, afu times *a, and tearing line, with a slope given by a linear fit to the experimentally determined J-*a data points. The 1981 procedure

yields lower Jic values (closer to the actual start of crack

growth) than the newer, 1987 standard method [49]. At 295

K, the J-*a data points in the blunting region did not

correlate to the standard slope for the blunting line.

Mills [50] observed similar behavior for austenitic stainless steels tested at room temperature and above. He found

that the data fit a blunting line with a steeper slope,

4aft*a. The steeper slope was used for analysis of specimens

tested at 295 K in this report because the steeper slope fit

the data better than the standard slope of 2aft*a. At 76 and

4 K, there was little difference in results for the two different slopes, so the standard slope was used for analysis.

120
82.6

z

Figure 11. (a) Drawing for flat, uniaxial tension specimen.

(b) Drawing for flat, biaxial tension specimen.

Blanks for toughness testing were either 26- or 13 -mm

thick, and were taken from the plate with a T-L orientation

so that the center of the plate is not tested, Fig. 12. Two

13 -mm thick blanks were machined with a L-T orientation.

Compact, C(T) specimens were designed according to the ASTM

121
standard, 12.7-mm thick for 4-K tests and 25-mm thick for

the other tests. Fatigue precracking was performed with an

R ratio of 0.1 at 20 Hz with a maximum K of 35 MPa m. The

fatigue precracking temperature was 76 K for 4- and 76-K

tests and 295 K for 295-K tests. Side grooves, 10% of the

thickness on each side, were machined in the 295- and 76-K

specimens after fatigue precracking. Specimens with T-L

orientation tested at 4 K were labeled with a prefix, A.

Specimens tested at 76 and 295 K were identified by AAA and

AA prefixes respectively.

The tests were conducted in stroke-control with a

servo-hydraulic controlled test machine, 100 kN load

capacity. The mechanical testing at cryogenic temperatures

requires a modified loading frame. The details of the exact

configurations for the tests described here are given

elsewhere [47,51].

3.3 Post-test observations

The load-displacement curves from the fracture

toughness tests were analyzed after the testing to determine

<

1

Rolling

Direction

Fig. 12. Relative location of 25 and 12.7-mm thick compact

specimens with T-L orientation in processed plate.

122
the value of 6 at J] C or at the end of the tests that were

stopped prior to Jio The value of <5 up to the point of Jj c can be calculated from the load-displacement curve for a

given specimen by the procedure described in BS 5762 [52].

The total <5 is the sum of the small-scale yielding (<S s) and

the large-scale-yielding ( <5 1 ) as shown in the equation,

The Vpi = plastic contribution to the mouth opening

displacement, W = specimen width, a0 = original crack

length, and k = rotational factor generally assumed to be

2.5. The elastic constants, Young's modulus (E) and

poisson's ratio (v)

,

are taken to be 200 GPa and 0.3 respectively, independent of test temperature [53].

After testing at each temperature, selected specimens

from uniaxial, biaxial, and triaxial stress states were

sectioned through the deformed region to observe

micromechanisms of ductile fracture. For the C(T) specimens

tested at 295 K, a wedge was placed in the crack to keep the

S constant during unloading and sectioning. One half of

each sectioned C(T) specimen was fatigue post-cracked to

reveal the extent of ductile crack growth on the fracture

surface while the other half was polished and etched for

metallographic examination. Three or four different

polished sections (serial sectioning) were observed for each

selected C(T) specimen. The distance between each serial

section was 1.25 to 1.5 mm.

The goal of the post-test observations is to observe

where void nucleation occurs and to characterize the strain

that exists at different locations in the specimen,

particularly the point where void nucleation starts. The

stress at any point can be estimated from the load and the

geometry. In uniaxial and biaxial stress states, the strain

is known and can be correlated to the change in microhardness (aDPH). In the triaxial stress state, the strain can

be inferred from *DPH, compared to *DPH of the tension

specimens strained to known values. The stress can be

estimated from FEM results. The maximum stress ahead of the

crack tip in a fracture mechanics specimen is a function of

distance from the crack tip and the applied energy. The

relationship developed by Anderson and Dodds [54] for smallscale yielding is used in this study.

where

and <5i = Vpt / l+k(a /W-a0)

(5)

(6)

(7)

123
4 . 0 RESULTS

4.1 Material Characterization

The grain structure of the austenitic matrix varied

depending on the location in the plate and was related to

the nonuniform structure found in the as-received bar.

Figure 13 shows the cross section of the annealed plate in a

schematic diagram that outlines three regions where the

microstructure varied. Region I is a fine-grained matrix,

austenite grain size between 40 and 50 /xm in diameter, with

a dispersion of spherical particles. A typical distribution

of the spherical particles is shown in a TEM micrograph,

Fig. 14. The majority of the particles are < 1 jxm in diameter and scattered randomly in the austenite. Based on the

chemical composition of the steel, the submicrometer particles are assumed to be V,Nb{C,N} precipitates. The precipitates will be referred to as microalloy particles. Region I

corresponds to the hatched region in Fig. 10 that was free

of sigma phase in the microstructure of the as-received bar.

Regions II and III of Fig. 13 correspond to the darker

cross hatched region in Fig. 10 that contained sigma in the

starting microstructure. Figure 15 shows the typical microstructure of Region II. The structure has a duplex grain

size with fine-grained areas like Region I plus areas of

extensive grain growth due to the hot deformation processing

Figure 13. Schematic drawing of the cross section of the

processed plate. Region I has a uniform fine

grain size. Region II has a duplex grain

size while Region III in the center of the

plate has a uniform large grain size.

124
Figure 14 . Representative TEM micrograph shows the distribution of submicrometer particles in the

processed plate after annealing. history

.

The small arrows in Fig. 15a point out the remanent of sigma

phase from the starting structure. The large arrow in Fig.

15a highlights an isolated nonmetallic inclusion. The

composition of the majority of inclusions included Mn, Si,

and S. The higher magnification view, Fig. 15b, shows

distinctive lines of the microalloy particles within the

coarse-grained regions (200 to 300 fim in diameter) . Region

III had a uniform, large grain size that was equal to the

coarse-grained areas in Region II.

SEM with the PRC program quantified the size and distribution of inclusions. The results for three separate

regions of transverse-oriented specimens taken from Region I

125
Fig. 15 Microstructure found in Region II of the plate

cross section. The large arrow shows a nonmetal

lie inclusion and the small arrows point out a

phase.

126
in the plate are shown in Table 3 . The area fraction of

inclusions varies from 0.01 to 0.05% of the total area

observed and is largely determined by the largest inclusions, > 3.8 nm in diameter, found in the distribution.

Most of the inclusions were spherical and < 1 /xm in diameter. The average inclusion spacing is the parameter that is

most representative of the complete inclusion distribution.

The spacing was calculated by taking the inverse square root

of the total number of inclusions divided by the total area

observed.

Table 3. Inclusion

PRC

distributions

program.

in 22-13-5 from

max. dia. (urn) Reaion 1 Reaion 2 Reaion 3

0.6 23 20 10

0.8 25 29 10

1.0 17 15 17

1.2 14 14 7

1.4 15 14 5

1.6 4 9 10

1.8 9 11 4

2.0 6 4 1

2.2 8 8 2

2.4 7 6 7

2.6 4 1 2

2.8 1 5 6

3.0 2 1 4

3.2 3 3 4

3.4 1 4 3

3.6 0 1 2

3.8 3 1 1

4.0 0 1 2

4.2 2 2 2

>4.t 2 i 1 4

# of inclusions 145 150 103

fields 30 36 28

Area fraction 0.0004 0.0001 0.0005

ave . spacing 95 nm 102 /Lim 109 nm

ave. diameter 1.4 nm 1.6 pm 1.9 nm

127
4.2 Mechanical testing

4.2.1 Uniaxial tensile test results

The test temperature has a large effect on the plastic

flow properties of 22-13-5. The complete load-vs. -time

curves to the point of fracture are shown in Fig. 16. At

295 K, the steel showed extensive nonlinear behavior before

the maximum load was reached (the region of uniform deformation) and relatively little between the maximum load and the

fracture load (the region of post-uniform deformation) . At

76 K, the steel had a higher load capacity (flow stress),

approximately twice that at 295 K. The shape of the load

vs. time curve at 76 K is essentially the same as at 295 K.

At 4 K, the initial flow stress is nearly three times higher

than 295 K, but the uniform deformation is interrupted at an

early point by load drops. The sawtooth behavior is caused

by local heating from plastic deformation. Each load drop

in the curve for the 4-K test reflects a local neck forming

in the gage section where the specimen was heated and then

strain hardened. The last load drop is associated with the

fracture of the specimen.

The true stress-vs . -true strain curves for representative specimens are shown in Fig. 17. The curves for 295 and

76 K stop at the point where the extensometer fell off the

specimen. The curve for 4 K stops at the point where serrated yielding began. The elastic modulus is nearly independent of test temperature, but the initial flow stress is

a strong function of temperature. Figure 18 shows the

strain hardening rates for each test shown in Fig. 17. A

decrease in test temperature produces a higher hardening

rate in these tests. The hardening rate at strains < 0.02

are a strong function of temperature and strain. Above 0.02

strain, the hardening rates become relatively constant with

strain, but still depend on temperature.

The summary of tensile properties from all the tests is

shown in Table 4. The specimens were taken randomly from

Regions I and II in the plate. The scatter in properties

may be attributed partially to the variation in the grain

size and density of microalloy particles found in Regions I

and II. The UTS and % EL values reported for the 4-K specimens are influenced by the heating of the specimen during

deformation. The fracture stress and the hydrostatic stress

concentration factor, oj/o0 (calculated according to [55]),

for the 4-K specimen were calculated based on the smallest

diameter present on the gage section away from the fracture.

128
T

N>i ‘QV01

129

Figure

16.

Uniaxial

tension

test

load

vs.

time

curves

for

22-13-5

as

a

function

of

test

temperature

showing

the

complete

load

history

up

to

fail-
2000 o

o

o

o

o

o

in o in

EdlAI ‘SS3H1S 3nni

O

<

130

Figure

17.

Uniaxial

tensile

test

true

stress

vs.

strain

curves

for

22-13-5

as

a

function

of

test
LO

o

d

<

CE

I

cn

g

!

cn

<

QQ) •P (0

10

u

c

•H

c

0

V

u

<0 £

c

•H

<c

•Pw

p

w

<D P

Q)

W

C

0) P

rH

<cX

<0 O

•H • C

C W 3

D > <w

00

(Vp

3

tr> *H

fa

Bd9 ‘3±Vd 9NIN3QH VH NlVdlS

131

plastic

strain

curves

for

22-13=5

as

tion

of

test

temperature.
Table 4. Uniaxial tension test results for

annealed 22-13-5 plate.

YS UTS El

SDec # (MPa) (MPa) _LI.1-

295 K

T-10 465 798 41

T-ll 438 793 44

T-l 400 779

T-3 430 788 41

T-4 398 752

76 K

T-8 1067 1430 39

T—

9

1034 1515 35

T-12 1030 1523 27

T-5 1097 1545 37

T-6 1101 1543

4 K

T-7 1272 1765 25

L-l 1320 1795 29

L-2 1448 1784 22

L-3 1425 1799 26

L-4 1439 1764 23

T-2 1264 1764

RA

(%)

fracture

stress Oj/On

1.25 2014 0.63

1.21 1780 0.63

1.32

1.06 1681 0.52

0.97 3026 0.48

0.56 2568 0.37

0.80 2988 0.45

0.88 3726

0.65 3117 0.42

T = transverse orientation and L = longitudinal orientation

4.2.2 Plane-strain tension test results (biaxial stress

state)

Decreasing the test temperature raised the flow

stresses in the plane-strain tension test, analogous to the

behavior in the uniaxial tension tests. Curves for representative load vs. change in specimen thickness (displacement) for plane-strain tension specimens at 76 and 295 K are

shown in Fig. 19. The specimen tested at 76 did not show a

load drop or necking prior to fracture. At a point near the

fracture load at 76 K in Fig. 19, the strain hardening rate

prior to fracture is greater than the flow stress. The

fracture was associated with a shear instability rather than

the plastic instability associated with necking. At a point

near the maximum load in the curve for the specimen tested

132
at 295 K, the strain hardening rate was almost equal to the

flow stress so that the steel meets the criterion for necking at the maximum load. Fracture is associated with the

local neck that forms after maximum load.

Table 5 shows the results of plane strain tension tests

at the three temperatures. The yield strengths of these

specimens do not correlate well with the uniaxial test

results in Table 4. The center of the specimen gage section

where the displacement was measured was located in Region

III of the plate (see Fig. 13) where the grain growth occurred during processing.

4.2.3 Fracture toughness test results (triaxial stress

state)

A typical P-u curve for a toughness test at 76 K is

shown in Fig. 20. The point where Ji c is measured is less

than the maximum load. For each increment of imposed displacement after Jic extra displacement and a load drop were

measured under fixed-grip conditions. The compliance appears to decrease during the load drops so that the crack

grows during the load drops. This behavior indicates that

the crack growth at cryogenic temperatures was not stable

even though it was ductile in nature. From the P-u curve in

Fig. 20, the J vs. *a is calculated from the area under the

curve and the change in compliance at each unloading point.

The data points are shown in Fig. 21 and were fit to a

blunting line with the standard slope and a tearing line.

The intersection of the two dashed lines is Ji»

Table 6 summarizes the fracture toughness test results

as a function of temperature. Figures 22, 23, and 24 show

the J-vs.-Aa curves for all the tests at each temperature,

while Fig. 25 summarizes the results as a function of temperature. Decreasing the test temperature results in lower

toughness. The three complete tests at 295 and 76 K, Figs.

22 and 23, were reproducible within the usual scatter expected in toughness testing. For the 4-K results (Fig. 24)

,

the first three complete tests were run and showed more

scatter. To check the first three tests, an additional five

specimens were tested. With the eight complete tests

together the results sort into two groups, HI-4K and LO-4K.

Two interrupted tests were also run at each temperature for

subsequent metallographic analysis of ductile fracture

micromechanisms. The C(T) specimens tested at 295 and 76 K

are physically larger than the 4-K specimens (refer to Fig.

12) and sample more material than the 4-K specimens because

133
n>i ‘avoi

134

Figure

19.

Plane

strain

tension

test

load

vs.

displacement

curves

for

22-13-5

at

295

and

76

K.
LOAD-LINE DISPLACEMENT, mm

Figure 20. Load vs. displacement curve for AAA-2 tested

at 76 K. The solid lines represent the imposed displacement, the dashed lines are for

the unloading compliance experiments, and the

dotted lines represent load relaxation under

fixed-grip conditions.

135
Table 5 Plane-strain

for annealed

tension test results

22-13-5 plate

Spec . # YS (MPa) UTS (MPa) Final Strain**

295 K

3 267 954 0.24

4* 293 967 0.21

8* 265 696 0.10

9* 304 958 0.22

12 307 801 0.12

76 K

2 1016 1893 0.18

5 720 2082 0.21

6* 952 1800 0.15

4 K

7 — 2040 0.08

10 1639 2015 0.03

11 ——

2187 0.09

* interrupted test, not broken

** fracture strain for failed specimen, or strain at the

point where the test was interrupted

136
800

O

E

E

LU

Q

3

U

rH <D

<0 >

U 5-1 3

<N U

I < 4J < c

< 0)g

5-1 0) 0 U

M-l (0

r-l 4J CU

0 W

rH *H

CLi'O

<d • « w

>

W T3

> «JO

hD «H

-P Q) W JZ

Q) -P

4-> g

w o

w p c

Q) «P -H

C

j: T3 c

o' a) 3

3 4-> o

o rd js

E-t h w

(N

<15 U

3

O'

•H

fa

Uiiu/N ‘iVdOElNI P

137

Fig.

20.
CRACK GROWTH, Aa (mm)

Figure 22. Toughness tests blunting and tearing lines

for 22-13-5 at 295 K, T-L orientation. Three

complete tests and two interrupted tests are

shown.

138
-

INTEGRAL

(N/mm)

1000

800 -

600 -

400 -

200

'0 0.5 1.0 1.5

CRACK GROWTH, Aa (mm)

Figure 23. Toughness test blunting and tearing lines for

22-13-5 at 76 K, T-L orientation. Three

complete tests and two interrupted tests are

shown

.

139
500

0.5 1.0 1.5

CRACK GROWTH, Aa (mm)

24. Toughness test blunting and tearing lines for

22-13-5 at 4 K, T-L orientation. Eight complete tests and two interrupted tests are

shown.

140
1000

0 0.5 1.0 1.5 2.0

CRACK GROWTH, Aa (mm)

Figure 25. Summary of all complete toughness tests for

22-13-5, T-L orientation.

141
Table 6 Fracture toughness results for 22-13-5

in T-L orientation

Specimen Jlc Kic(J) Tearing Modulus Si

4_K

A-2 99 N/mm 141 MPa m 23 46 /im

A3122 143 14 44

A4

59 109 7 28

A-5 58 108 6 28

A-6 62 111 8 26

A-7 58 108 5 28

A-8 110 148 18 46

A-9 112 150 18 49

A-10* 94 137 — 36

A-ll* 51 101 —

2

76 ..K AAA-1 267 231 28 156

AAA-2 330 257 26 156

AAA3220 210 30 131

AAA-4* 187 193 — 97

AAA-5* 54 104 — 33

295 K

AA-1 524 324 222 280**

AA-2* 228 213 — 220**

AA-3 540 329 197 385

AA-4* 100 100 — 79**

AA-5 576 339 198 387

* interrupted test

** measured from specimen

the plastic zone size is larger at higher temperatures. The

toughness specimens were machined so that the crack path in

the specimens would avoid the center region of the plate,

Region III of Fig. 13, where extensive grain growth occurred • Two toughness tests were run at 4 K on specimens with

L-T rather than T-L orientation. The J vs. *a curves for

the two L-T tests and one of the HI-4K and one of the L0-4K

tests are shown in Fig. 26. The data points for the L-T

specimens fall in between the HI-4K and L0-4K results. The

142
toughness is not a strong function of specimen orientation

because the matrix material is relatively isotropic (little

difference in tensile properties in the two orientations)

and the second-phase particles in the matrix are small,

widely spaced, and mostly spherical (Table 3)

.

4.3 Post-test observations from uniaxial stress state

4.3.1 Room temperature tests

Fracture at 295 K was a cup-cone type of failure where

there is a flat central region surrounded by a ring of shear

lips, typical of ductile fracture caused by MVC. The fracture surface from a transverse tension specimen shown in

Fig. 27 is typical. The low magnification view, Fig. 27a,

shows the dimpled rupture features of MVC in the center of

the specimen contained within an annular ring of shear lips.

The biggest features in the central region are elongated

troughs with rough surfaces characteristic of extensive

plastic deformation. The troughs form at isolated nucleation sites and grow together by plastic deformation. The

elongated troughs are surrounded by smaller, equiaxed

dimples with a smooth surface. The dimples form from closely spaced nucleation sites that coalesce with little associated plasticity.

A cross section through a failed specimen oriented

across the rolling direction and tested at 295 K is shown in

Fig. 28. The profile of the fracture surface is shown at

the top of the figure. The dark features in the unetched

matrix represent voids. The first voids to nucleate are

scattered through the necked region of the specimen, indicating that the strain for nucleation is greater than the

uniform strain. The largest voids are elongated along the

tensile axis, which is transverse to the rolling direction

of the plate from which this specimen was taken.

The primary void nucleation sites in the uniaxial

stress state at 295 K are the sigma phase left from the asreceived structure. The etched microstructure below the

fracture surface of the specimen from Figs. 27 and 28 is

shown in Fig. 29. Sigma is the light etching phase

elongated across the figure. The dark patches within the

sigma phase are voids that have nucleated by particle cracking. Void nucleation was limited to the elongated sigma.

The troughs observed on the fracture surface correspond to

the spacing of sigma in the microstructure.

143
400

CM

LO pi i 0) 1 -H

- o P- £ Eh 0

• 4-> 0) T— a

L-T

1

L-T 2

T-L

A-8

T-L

A-4

for th K

s

- O

•

•H O

W £ »T

4J 1 C TJ O

•H Q) J

0 P

< O • d flJ (J

a i in <a e eh -P 0

- _ CM <0 0 <0 T— E « TJ

o E

<a c

^ Xf (C O

o CO

• -p —

W rtJ CO > 1 0 < U1 <

b

h- <D

4-> g C

W *H O

o in UJ Q> 0 g

4-> <D *H

- h* Q a o • MB 1)^

A o W d xr * 0) Eh W 1 *3* G 1 <

JS PI X ~

II o O' xr _ H in 0 0 1 c

0 5 H o

o

• o

H4JK g

o

• D < °o**

in

CM

•

• VO

CM <3 ^ o

<DP

1 L_

0

i i 00

O'

•H

Cx. o o o o o o

in o in o in o

CO CO CM CM 'rm

_ £ _

% lUUJ/N ‘1VUD31NI P

144
toughness is not a strong function of specimen orientation

because the matrix material is relatively isotropic (little

difference in tensile properties in the two orientations)

and the second-phase particles in the matrix are small,

widely spaced, and mostly spherical (Table 3)

.

4 . 3 Post-test observations from uniaxial stress state

4.3.1 Room temperature tests

Fracture at 295 K was a cup-cone type of failure where

there is a flat central region surrounded by a ring of shear

lips, typical of ductile fracture caused by MVC. The fracture surface from a transverse tension specimen shown in

Fig. 27 is typical. The low magnification view, Fig. 27a,

shows the dimpled rupture features of MVC in the center of

the specimen contained within an annular ring of shear lips.

The biggest features in the central region are elongated

troughs with rough surfaces characteristic of extensive

plastic deformation. The troughs form at isolated nucleation sites and grow together by plastic deformation. The

elongated troughs are surrounded by smaller, equiaxed

dimples with a smooth surface. The dimples form from closely spaced nucleation sites that coalesce with little associated plasticity.

A cross section through a failed specimen oriented

across the rolling direction and tested at 295 K is shown in

Fig. 28. The profile of the fracture surface is shown at

the top of the figure. The dark features in the unetched

matrix represent voids. The first voids to nucleate are

scattered through the necked region of the specimen, indicating that the strain for nucleation is greater than the

uniform strain. The largest voids are elongated along the

tensile axis, which is transverse to the rolling direction

of the plate from which this specimen was taken.

The primary void nucleation sites in the uniaxial

stress state at 295 K are the sigma phase left from the asreceived structure. The etched microstructure below the

fracture surface of the specimen from Figs. 27 and 28 is

shown in Fig. 29. Sigma is the light etching phase

elongated across the figure. The dark patches within the

sigma phase are voids that have nucleated by particle cracking. Void nucleation was limited to the elongated sigma.

The troughs observed on the fracture surface correspond to

the spacing of sigma in the microstructure.

145
Figure 27. The fracture surface appearance of tension

specimen with transverse orientation with

respect to the rolling direction, tested at

295 K.

146
147

Figure

28.

Cross

section,

unetched,

through

transv

uniaxial

tension

specimen

broken

at

295
Figure 29. Microstructure below the fracture surface of

tensile specimen cross section shown in Fig.

28. The darker patches across the center of

the picture are voids.

148
The micromechanisms of ductile fracture at 295 K are

interpreted from the observations as follows. Voids form at

sigma phase after large strains at a stress near the UTS.

Voids grow in the ligaments between cracked particles to

form the elongated voids observed in the cross section below

the fracture surface. At the fracture stress, the equiaxed

dimples nucleate from microalloy particles in the matrix

between the elongated voids. The equiaxed dimples are

assumed to be associated with the submicrometer size microalloy particles even though the particles are not observed

on the fracture surface. Fracture surfaces of martensitic

steel [56,57] display submicrometer, equiaxed features like

those found on the fracture surfaces of 22-13-5. For the

martensitic steels, carbon extraction replicas proved unambiguously that the equiaxed dimples nucleate at undissolved

carbides in the steel.

4.3.2 Liquid nitrogen tests

When the test temperature was reduced from 295 to 76 K,

the macroscopic fracture changed from a cup-cone fracture

associated with MVC in the center of the specimen to a

mixture of MVC and shear. A low magnification view of a

failed 76-K specimen is shown in Fig. 30. The specimen

broke into three pieces where the necked region separated

Figure 30. Macrograph of uniaxial tension specimen broken at 76 K where the necked region has separated intact from the ends of the tension

specimen.

149
cleanly from the ends of the specimen, leaving a hole in the

center of the necked piece. The fracture surface is all

shear lip.

The fracture surface of a second tension specimen

tested at 76 K is shown in Fig. 31. Most of the surface

consists of shear lips. At higher magnification (Fig. 31b)

,

the center of the fracture surface shows large facets, 20-50

jLim in diameter, surrounded by smaller dimples, 1-2 /im in

diameter, that lie on planes inclined to the tensile axis.

A polished section taken from the tension specimen

tested at 76 K is shown in Fig. 32. The profile of the

fracture surface is made up of a series of intersecting

shear planes with a small, 150 nm long, flat section in the

center. Voids are observed below the fracture surface, but

less elongated than observed in the cross section of the

295-K specimen (Fig. 28)

.

At 76 K, the micromechanisms of ductile fracture are

different than those at 295 K. Shear fractures cover a

larger percent of the surface. The primary void nucleation

sites are still the patches of sigma left in the microstructure after thermomechanical processing. After void nucleation, strain is localized into shear bands between voids.

Within the bands, a smaller set of voids are nucleated as a

type of void sheet coalescence. In an extreme case like

that observed in Fig. 30, the fracture surface was all shear

lip because strain became localized into shear bands between

a single void near the center of the specimen and the outer

surface of the specimen.

4.3.3 Liquid helium tests

The appearance of the fracture surfaces of specimens

tested at 4 K is similar to the observations made at 76 K.

Away from the fracture surface of 4-K specimens, several

diffuse necks are observed along the gage length of the

specimen, as well as at the point of fracture. The specimen

heating occurs locally along the gage section causing extensive plastic deformation, producing neck formation, and

strain hardening the matrix. Subsequent strain occurs away

from the neck because the matrix is softer. The temperature

can rise as much as 100 K in the neck during testing [47].

Given the local heating of the 4-K specimens, it is not

surprising that the fracture surfaces of 4-K specimens

resemble those at 76 K.

150
Figure 31. The fracture surface appearance of transverse

tension specimen broken at 76 K.

151
Figure 32. The cross section, unetched, through transverse uniaxial tension specimen, broken at 76

K.

152
Figure 33. The cross section through biaxial tension

specimen strained at 295 K, but not broken:

(a) center of gauge section and (b) edge of

gage section. Arrow shows elongated second

phase stringer.

153
Figure 34. Microstructure found in specimen shown in

Fig. 33: (a) center of gage section where a

patch of o phase is cracked, and (b) edge of

gage section.

154
Figure 35. The fracture surface appearance of biaxial

tension specimen broken at 295 K.

155
Figure 36. The cross section, unetched, through biaxial

tension specimen broken at 4 K: (a) center

of gage section, and (b) edge of gage section.

156
c

0)g 0) *h .c u -P

<da c

in o

<4-1

w o 0)u

0) 3 u -P C <0

(0 0 P <4-1 10

0)u

a>

as • (0 <D O' o

a) c <o U *H <4-1 (0 £ p <4-10 3 u sz in 3 in

in a) * p

<D 3 P -P 3 O

•P -P <0 U d P

(0 <4-1 P T3

<4-1 Q) P

•P

d

fl) 10 <D £ o) s: H «p in

r>©p3

O'

*H

157
200

<0

(0

(HdQ) SS3NQUVH Nl 30NVH0

O

CM

i

<

cch

(0

o

hcn

<

Cl

158

Figure

38.

The

change

in

microhardness

(adph)

as

a

function

of

strain

in

uniaxial

and

biaxial

tensile

tests

at

4,

76,

and

298

K.
Figure 39. Deformation substructure found in specimen

deformed at 295 K: (a) after 0.02 strain, and

(b) after 0.05 strain.

159 i
Figure 40. Deformation substructure found in specimen

deformed at 76 K; (a) after 0.03 strain, and

(b) after 0.12 strain.

160
associated fracture surface will be examined first. Finally, the adPH around the crack tip of each specimen is shown.

4.7.1 Room temperature tests:

The test of specimen AA-4 was stopped at a point where

the section was just starting to yield (transition from

small-scale yielding to large-scale yielding); see Fig. 22.

A representative section through the blunted fatigue crack

is shown in Fig. 41. 6 at this point is about 0.08 mm. The

grain size next to the crack tip in this section is larger

than the bulk material. The crack at this point was located

in Region II of the plate microstructure? see Fig. 13.

Along the crack front, the grain size varies between 0.04

and 0.2 mm. There is a cluster of second-phase particles,

probably sigma phase, ahead of the crack tip, but the distance from the tip is greater than 26 , the process zone

size.

The next set of observations come from specimen AA-2,

interrupted at a point just below Ji c (Fig. 22) . The profile

of the blunted crack tip is shown in Fig. 42. <5 at this

point is 0.22 mm. The blunted tip has a rough, irregular

shape and a crack has grown about 0.05 mm ahead of the

blunted tip. A void is observed at an angle below the crack

path, 0.25 mm or about a distance 6 away from the crack

present at the blunted tip.

The fracture surface from the second half of specimen

AA-2 is shown in Fig. 43. The region created during the

interrupted test is a narrow ribbon of material across the

figure, similar to the SZ discussed in the introduction. At

the higher magnification (Fig. 43b) , the SZ appears as a lip

sticking up above the fatigue features that formed during

precracking and post-cracking. The surface of the SZ looks

like a series of slip steps. The fatigue features above

(post-cracking) and below (pre-cracking) the SZ are the

same.

Specimen AA-1 was tested to a crack growth of about 1.8

mm. A cross section through the crack is shown in Fig. 44.

6 for initiation was estimated from this picture to be about

0.28 mm by subtracting the 6 at the tip of the original

fatigue crack from the 6 at the left edge of the figure.

The actual crack tip after 1.8 mm of growth remains sharp.

Voids have nucleated and grown above and below the crack at

elongated, second-phase particles in the matrix. The

161
Figure 41. The cross section through compact specimen

AA~4 tested at 295 K. Crack growth direction

is left to right.

162
f

.

V

r ^ *\

r - -

r Nl

i

Figure 42. The cross section through compact specimen

AA-2 tested at 295 K. Crack growth direction

is left to right.

163
Figure 43. The fracture surface appearance of specimen

AA-2 tested at 295 K. The SZ formed during

the toughness test is marked with arrows.

164
co <4-1 •H O

G P (DOC g 0) O

•H P *H O -H P

Q) T3 «J a o w .c o P H P 2 U OT3

(T3 P <D CU O'P e o o * c o o

(0 .c p

C'CJ 3op • ,C W P

in c ON O CN •HP +j O (0 0) *00pp< o3po

<dp

a.

a)

Ul 730) in P

in in o CD ppp o o H iH

Q> I .C < Wp <D G 3

O' O'

•H *H P -P3 O 4-1 P

i-“i

p (T3 cH

O'HP op3

O'

•H

fa

165
initial crack growth was directly ahead of the fatigue crack

where the stress is a maximum. After about 1 mm of crack

growth, the crack moved out of the plane of maximum stress,

apparently to link up with neighboring voids that had formed

ahead of the continuous crack tip.

The fracture surface from the second half of specimen

AA-1 is shown in Fig. 45. The fatigue precrack features

(the bottom of Fig. 47a) lead into the SZ, followed by the

dimpled rupture features characteristic of MVC (at the top

of Fig. 47a) . At a higher magnification in Fig. 47b, the

dimples show a variety of sizes and shapes. The larger,

elongated dimples resemble the troughs observed on the

fracture surface of uniaxial tension specimens tested at 295

K, see Fig. 27, which formed from the sigma phase. Between

the elongated dimples, equiaxed dimples with varying size

cover the rest of the surface. The equiaxed dimples nucleate at the submicrometer-size microalloy particles.

The adPH around the crack tip of the three 295-K C(T)

specimens examined above are shown in Fig. 46. The maximum

value of adph is normally found next to the crack tip in

each specimen and the maximum value does not change appreciably with orientation around the crack. The gradient in

*DPH is sharper ahead of the crack than perpendicular to the

crack plane. The *DPH decreases to a value between 20 and

40 at a distance of 2 to 3 times 6 ahead of the crack tip.

Perpendicular to the crack, the comparable *DPH are at a

distance of 10 to 20 times 6 . The C (T) specimens from the interrupted tests at 295 K

were wedged open so that the actual values of 6 could be

measured from the serial sectioning experiments. The measured values of 6 for the interrupted tests were used to

derive a rotation factor that would give the same values of

6 calculated from Eq. 5. Table 7 shows the results of these

calculations. For specimen AA-2 (interrupted near Ji<), the

calculate factor was 2.6, nearly equal to the standard value

of 2.5. However, for the specimen interrupted at the lowest

applied energy, AA-4, the calculated factor was 4.2.

The dominant micromechanism of crack growth in the

triaxial stress state at 295 K is MVC. The plastic strains

defined by DPH measurements associated with the ductile

crack growth are distributed over large distances around the

crack tip, 10 to 20 times 6 , but the micromechanisms of

crack growth are limited to a region directly ahead of and

close to the crack tip, 1 to 3 times 6 . The observa166
Figure 45. The fracture surface of compact specimen AA-1

tested at 295 K. The SZ is marked with arrows in the center of the part (a)

.

167
200 CO

oCNJ o

CO

o

m

cvi

in

E

E

(HdQ) SS3NQUVH Nl 30NVH0

I 3

c

0) g

•H

o

Q)a, (0

4-1 0

a

•H

+J*

O

(0

J-) O

d) x: • -P «

TJ ID C <T\ 3 CM O

U -P

(0 (0 ai a) o -P H W

0) d) 4J

jc

Eh

<0 VO

a)u

3

CP

•H

fc,

168
200

in

E

E

VN HI

o

<

<n

a

!

I

5

(HdQ) SS3NQUVH Nl 30NVH0

169

Figure

46b.

The

adph

around

the

crack

tip

of

specimen

2

tested

at

295

K.
200

w

LU

O

Z

<H

CO

Q

5

(HdO) SS3NQHVH Nl 3DNVH0

170

Figure

46c.

The

aDPH

around

the

crack

tip

of

specimen

1

tested

at

295

K.
Table 7 The 6 calculations for

tested at 295 K

compact specimens

SDec. # 5 (mm) Ls Ipill |o 1 1

!

Bp(adi)

AA-2 0.22 0.06 0.16 2.6

AA-4 0.08 0.027 0.054 4.2

tions on the fracture surface (Fig. 45b) show that the first

crack to grow from the fatigue precrack is associated with

slip steps and not with MVC. The slip steps are characteristic of shear localization on the scale of the deformation bands observed in TEM (Fig. 39a) . The microalloy particles are scattered randomly throughout the matrix with a

spacing of 1 to 5 urn so there are numerous potential nucleation sites inside the PZ at 295 K, PZ = 80 yin for AA-4 and

440 ym for AA-2 . The stress required to nucleate voids from

the microalloy particles is higher than the local stress in

the two interrupted tests, AA-4 and AA-2. At a higher

applied energy and 6 that must be approximately equal to the

average J

lc

value for the room temperature tests, the local

stress exceeds the value of o

c for the microalloy particles.

The value of o

c for nucleation from sigma phase is lower

than that for nucleation at the microalloy particles, but

the distribution of the weaker sigma particles is such that

the probability of finding a particle of sigma in the PZ is

small

.

4.7.2 Liquid nitrogen tests

For the C (T) specimens tested at cryogenic temperature,

a wedge could not be used to keep the crack tip open after

unloading and sample preparation. The elastic component of

6 is lost on unloading and the observed 6 is less than the

actual value of 6 during the test. The value of 6 for all

low-temperature C(T) specimens was calculated from Eq. 7

with a rotational factor of 3.0, a compromise value inbetween the two values calculated from the 295-K results.

At the next lower temperature (76 K) , specimen AAA-5

was interrupted at the transition from small-scale yielding

to large-scale yielding; see Fig. 23. The polished and

etched profile of the crack-tip region is shown in Fig. 47.

171
The calculated value of 6 was 20 nm and compared favorably

to the measured value of 6 from Fig. 47. The fatigue precracked surface profile shows slight undulations and occasional branching, suggesting a crystallographic nature to

fatigue crack growth. The crack tip does not appear rounded, but blunted to a sharp corner.

At a higher applied energy, specimen AAA-4 was interrupted near J Ic

. One profile through the blunted, crack-tip

region is shown in Fig. 48. The calculated value of 6 has

increased to 74 jim. A short crack is present at this point

in the toughness test. The new crack that has formed ahead

of the blunted fatigue crack is sharp with a value of 6 of

about 5 jim. Another section from AAA 4 is shown in Fig. 49.

Voids, 10 to 2 0 iim in diameter, are associated with the

short crack in this section. Higher OPH values were associated with regions of cracking and void growth like Fig. 49

compared to the OPH found in Fig. 48.

The fracture surface appearance of the other half of

AAA-4 is shown in Fig. 50. At the lower magnification (Fig.

50a) , a SZ is difficult to distinguish from the fatigue preand post-cracking features. In an isolated region along the

crack front, a large, 20 nm, void has opened up with smaller, 2 to 5 \im , dimples surrounding the cavity. Presumably,

this is a result of the nonuniform distribution of sigma

phase in the matrix.

The profile from a completed test, AAA-2, in which the

average crack growth was 2.0 mm, is shown in Fig. 51. The

location of the original fatigue precrack is shown by the

arrow. The apparent from this picture is 0.12 mm. The

6i calculated from equation 7 is 156 \im. The crack growth

direction is predominantly straight ahead of the crack tip,

but does take discrete jumps out of the plane of maximum

stress. The crack tip remains relatively sharp at the end

of crack growth with a 6 of about 15 |im.

The fracture surface from the second half of C(T) AAA-2

is shown in Fig. 52. The dimples are more or less spherical

in shape, but the size varies from submicrometer to 25-30 iim

in diameter. The large, elongated dimples are present, but

are not as prominent as those found in Fig. 45 for the

fracture surface of the C(T) tested at 295 K.

The deformed zones measured around the crack tip in the

76-K specimens are shown in Fig. 53. Again, the maximum

OPH directly ahead of the crack tip is similar to the

172
Figure 47. The cross section through compact specimen

AAA-5 tested at 76 K. Crack growth direction

is left to right.

173
Figure 48. The cross section through compact specimen

AAA-4 tested at 76 K. Crack growth direction

is left to right.

174
Figure 49. A second cross section through compact specimen AAA-4 tested at 76 K.

175
Figure 50. The fracture surface appearance of compact

specimen AAA-4 tested at 76 K. Crack growth

direction is bottom to top.

176
co

•H

C 4-> <D 03 g 0

•H 0 0 .H 0

Q4T3 w 0

-P P 0 o c

03 0

0,^3 g0 5 u 0u A U

tP< 30u • a £ « ‘H P P

VO c r- ^ 0 0

*H P 03 P 03 P 0 O 0 T3

(/3 0r-! P 03

C/3 C/3C

C/3 0 -H

04J S' P -H 0 CN P I o 0 < H

in0u3

cn •H

fci

177
Figure 52 . The fracture surface appearance of specimen

AAA-2 tested at 76 K. Crack growth direction

is bottom to top.

178
200

in

in

ID CM co

(HdQ) SS3NQUVH Nl 30NVH0

179

Figure

53a.

The

adPH

around

the

crack

tip

of

specimen

AAA-5

tested

at

76

K.
200

in

o

in

CD CO *3*

(HdQ) SS3N0UVH Nl 30NVH0

180

Figure

53b.

The

adPH

around

the

crack

tip

of

specimen

AAA-4

tested

at

76

K.
200

in

o

in

o

CM

(HdQ) SS3NQUVH Nl 30NVH0

181

Figure

53c.

The

*DPH

around

the

crack

tip

of

specimen

AAA-2

tested

at

76

K.
maximum value perpendicular to the tip, but the deformed

zone is larger perpendicular to the crack. Directly ahead

of the crack, the deformed zone is approximately twice the

value of <5 for each respective specimen. The deformed zone

directly ahead of the tip in specimen AAA-2 (Fig. 53c) does

not show the smooth exponential decrease observed in other

specimens.

The micromechanisms of fracture in toughness testing at

76 K are different than those observed at 295 K. No SZ was

observed between the fatigue precracked and post-cracked

regions. The dominant features associated with the toughness test on the fracture surface are small, 2 to 5 nm

dimples nucleated from microalloy particles. The local

stress due to the applied energy for AAA-5 was less than oc for void nucleation from microalloy particles. With the

higher applied energy for specimen AAA-4, the local stresses

were high enough to form a crack close to the tip in the PZ.

Isolated voids from microalloy particles are not visible

ahead of the crack tip. Once the critical stress for void

nucleation was reached, the voids formed a crack without

additional applied energy. However, the stresses at the

edge of the PZ were not high enough to nucleate voids. With

higher applied energy, taken to be equal to the average Jj c value for AAA-1, AAA-2, and AAA-3, the local stresses ahead

of the PZ exceeded ac for nucleation from the microalloy

particles and the crack advanced through the PZ.

4.3.10 C (T) specimens at 4 K

The specimen interrupted at the lowest load is specimen

A-ll , see Fig. 24. The applied energy at this point is less

than Jic for any of the completed tests. The crack-tip

profile is shown in Fig. 54. The crack tip has a very small

value of 6, about 7 jim measured from the picture. The

calculated value of 6 was 2 /m. No evidence of damage was

observed at magnifications at high as lOOOx.

Specimen A-10 was tested to an applied energy above Ji c for LO-4K specimens, but less than Jic for the HI-4K specimens? see Fig. 24. The profile from the blunted crack tip

of A-ll is shown in Fig. 55. The crack has blunted to a

square shape. The calculated 6 was 36 nm. The 6 observed

in the picture is similar to that observed for A-ll in Fig.

54. A maximum of 0.11 mm crack growth was observed ahead of

the crack tip. The crack growth is directly ahead of the

fatigue precrack.

182
Figure 55. The cross section through compact specimen A

10 tested at 4 K with T-L orientation.

183
Figure 54 . The cross section through compact specimen A11 tested at 4 K with T-L orientation.

184
The fracture surface from the second half of A-10 is

shown in Fig. 56. A thin, discontinuous ribbon of material

created during the test is visible between the fatigue

precrack on the bottom and the post-crack on the top. At

higher magnification (Fig. 56b) , the fracture surface

created during toughness testing is made up of small dimples, 2 to 5 /Ltm diameter. The fatigue precrack just below

the MVC features exhibits a dispersion of small particles on

the fatigue features.

A specimen tested to a crack growth of about 2 mm from

the LO-4K test group was sectioned for observation. The

crack-tip region from specimen A-4 is shown in Fig. 57. The

approximate position of the original crack-tip is indicated

by the arrow. The profile of the growing crack is similar

to the appearance of fatigue crack growth profile.

The scatter measured in the 4-K specimens during testing was not attributable to differences on the fracture

surface of the specimens. A representative area from the

fracture surface of a L0-4K specimen is shown in Fig. 58.

The typical appearance of the fracture surface of a HI-4K

specimen is shown in Fig. 59. A SZ is not visible in either

Figs. 58 or 59 between the fatigue precrack and dimpled

rupture features of MVC. The dimpled rupture features for

Figs. 58 and 59 are similar. The larger, 20 to 30 nm dimples are elongated in the crack growth direction and are

surrounded by smaller, 1 to 2 /m, equiaxed dimples. There

were no systematic differences in fracture appearance between T-L specimens that could be related to Ji c at 4 K.

The fracture surface of the L-T specimens tested at 4 K

were significantly different from the surfaces of the T-L

specimens. Figure 60 shows the fracture surface appearance

of specimen L-T 2. Dimples as large as 50 pm in diameter

are observed surrounded by smaller, 2 to 20 diameter

dimples. The overall features have a more equiaxed appearance than those found in Figs. 58 and 59 for the T-L specimens. The difference in fracture appearance between T-L and

L-T specimens at 4 K was not reflected by the toughness; see

Fig. 26.

The deformed zones around the crack tip of specimens A10, A- 11 , and A-4 are shown in Fig. 61. The maximum value

of aDPH around the crack tip is about 80, smaller than the

maximum value of 160 to 200 measured at 76 K (Fig. 53)

.

At

4 K, the deformed zone ahead of the crack tip is spread over

185
Figure 56. The fracture surface of compact specimen A-10

tested at 4 K with T-L orientation. Arrows

mark the original fatigue crack tip.

186
I • < CH

0 <U C'H C

Q) .p *H

6 03 O'

•H *H

o c ^

<D 0) O

CU*H

W U <W O 0

+J

„ U J c

03 I O

a E-* *H

g 4J

O 42 03 O -P O

•H O

42 5 <—

I

O'

3 * d) 0 p

3-1 ^ 03 42 g P -P •H

03 X

C 0 • 0 T3 P X

•H 0) a 0

-P -P a ns o in 03 p

0 0) 0

W -P in 0) 3 p

in ^

0 a

in « 42

0 in 0) p I 3

0 o 5 O'

PI 0 •H

Q) W ^ 4-> 42 P <0 &n ^ < *w

r*

in

o

P

3

O'

•H

fa

187
Figure 58. The fracture surface appearance of a L0-4K

compact specimen with T-L orientation. Crack

growth direction is bottom to top.

188
Figure 59. The fracture surface appearance of a HI-4K

compact specimen with T-L orientation. Crack

growth direction is bottom to top.

189
Figure 60. The fracture surface appearance of a compact

specimen tested at 4 K in the L-T orientation.

1 90
100

I

CL

O

CD

CM

CM

COW LO o i II ^

T5 rII II II

|

n

. o o

<x> <x> CD

U)(f)

o

CO CD

< O

£

* E

*3- CO _• O• r- o

II

I <

to

in

o

O

CO

o

CM

CO

o

CO

o o

CM

©3

E

E

Hi

O

(0

o

I

(HdQ) SS3NQaVH Nl 30NVH0

191

Figure

61a.

The

aDPH

around

the

crack

tip

for

specimen

11

tested

at

4

K

with

T-L

orientation.
100

Q.Q

CO

CO

CM

II o>

T3 0>

i

• s>°

o

CO GQ

E

CM

CO 09©

<"

oo

CM

CM

o i

II II II

<x> <x> ©

< O

<

in

o

9 ©

CO

o

CM

E

E

LU

o

w

Q

co

o

CD

o o

CM

I <

(HdQ) SS3NQUVH Nl I30NVH0

192

Figure

61b.

The

aDPH

around

the

crack

tip

for

specimen

10

tested

at

4

K

with

T-L

orientation.
100 GO CD CNI

ir> o

o

00

o

CM

o

©

E

E

l <

(HdQ) SS3NQHVH Nl 3DNVH0

193

Figure

61c.

The

*DPH

around

the

crack

tip

for

specimen

4

tested

at

4

K

with

T-L

orientation.
a distance larger than 26. At any given distance away from

the crack, the aDPH is greater at 0 = ± w/2 than at 0 * 0.

The micromechanisms of fracture at 4 K are similar to

those found at 76 K. The local stresses in specimen A-ll

were less than ac for void nucleation from the microalloy

particles. With a slightly higher applied energy, A-10, the

local stresses were greater than ac close to the crack tip

in the PZ, but not ahead of the PZ. At the applied energy

equal to Ji c for the HI-4K specimens, the local stresses are

equal to ac ahead of the PZ.

For the 4-K specimens tested with a L-T orientation,

the process appears to be the same as that found for the T-L

orientation. The orientation has no influence on the matrix

flow properties and does not change the relative shape of

the spherical, microalloy particles with respect to the

loading direction. The sigma phase particles are not spherical, but are pancake-shaped and the relative shape will be

different for the L-T orientation. The lack of elongated

features on the fracture surface of the L-T specimens suggest that the interface between the matrix and sigma particles is stronger in that direction.

5.0 DISCUSSION

The variations in YS and toughness as functions of

temperature reported here for 22-13-5 are typical of other

austenitic stainless steels [37-41]. The main question is,

what are the basic mechanisms that control toughness? The

metallographic observations of deformed and broken specimens

are not particularly different from what was previously

observed in ferritic steels [12-15, 30-32]. What makes this

work different is the emphasis on void nucleation, or more

correctly, a lack of void growth. No special techniques

have been used to measure the point where voids first separate from the matrix. Instead, this work concentrates on

the global conditions required to nucleate and grow voids to

a size that can be observed with light microscopy.

In this discussion, the observations with respect to

the different micromechanisms of ductile fracture are reviewed. The second part of the discussion concentrates on

quantifying the conditions required for significant void

nucleation as a function of test temperature and stress

state. Models for toughness are discussed in light of the

observations of micromechanisms in 22-13-5. The effect of

194
test temperature and stress state on the behavior of the

material are considered last.

5.1 Micromechanisms

The results of this study are summarized in terms of

the micromechanisms of ductile fracture in Table 8. When

two micromechanisms are listed, the first is dominant. At

the highest test temperature (295 K) , MVC was the dominant

micromechanism observed. Void nucleation occurred at the

sigma phase particles before void nucleation occurred at

microalloy particles. At lower test temperatures in the

uniaxial and biaxial stress states, shear fracture due to

strain localization became more important. Strain localization interrupted the uniform deformation in the matrix

that leads to MVC. For shear fracture to be the dominant

micromechanism, no significant void nucleation can be present.

In the triaxial stress state, void nucleation from

microalloy particles is the most important micromechanism

related to toughness. Shear fracture manifested by the SZ

was visible only at room temperature. Strain localization

without prior void nucleation is an important metallurgical

phenomenon, particularly in metal forming operations [6,

57]. The conditions existing in the tests where shear

fracture occurs represent a subcritical point for the nucleation of a significant volume fraction of voids. Because

shear fracture is not significant in toughness testing of

22-13-5, the discussion concentrates on MVC, the micromechanism observed in toughness testing.

5.2 Void nucleation conditions

The data presented in the results section emphasize the

relationship between void nucleation and crack formation in

22-13-5. Two different approaches to quantify the stress

required to nucleate voids were discussed in Section 1.0.

The data for uniaxial and biaxial stress states were analyzed by these two criteria, and the results are summarized

in Table 9 for the micromechanisms that lead to failure.

The specimen at 4 K in uniaxial stress state does not

actually fracture at 4 K because the specimen heats locally

during fracture [47], so the value of ac here refers to the

maximum stress on the minimum cross section that stayed at 4

K. The critical stress for void nucleation from the large

195
Table 8. Summary of types of ductile fracture in

22-13-5.

at 295 K

Uniaxial stress

state

Biaxial stress

state

Triaxial stress

state

at 76 K

Uniaxial stress

state

Biaxial stress

state

Triaxial stress

at 4 K

Uniaxial stress

state

Biaxial stress

state

Micromechanisms Comments

MVC

MVC +

shear

MVC +

shear

voids nucleate from

sigma phase at UTS,

smaller voids nucleate at microalloy

particles at fracture stress

splits form from

sigma phase at UTS,

remaining ligament

fails by shear

SZ forms by shear,

voids form

at microalloy particles

shear +

MVC

shear

strain localization in neck after

UTS, voids nucleate

from sigma phase

strain localization before UTS

MVC voids form at

microalloy particles

shear +

MVC

shear

strain localization

in neck before UTS,

voids nucleate from

sigma phase

strain localization before UTS

196
Table 8. Summary of types of ductile fracture in

22-13-5 (continued)

Micro

mechanisms Comments

Triaxial stress MVC

state

voids form at microalloy particles

second-phase particles at 4 K in a uniaxial stress state

must be greater than the value shown in Table 9.

For the biaxial stress state, the ac calculated for

void nucleation from sigma at 295 K is approximately the

same as the ac for nucleation from sigma at 295 K in uniaxial stress state. The value of ac calculated for the

biaxial stress state at cryogenic temperatures (shear fracture) in Table 9 is less than the value required to nucleate

a significant volume fraction of voids necessary for MVC.

Criterion 2 generally predicts a higher value than criterion 1? the difference is larger for uniaxial than the

biaxial stress state. The most striking point is that the

values are higher at cryogenic temperatures than at 295 K.

The analysis shows that void nucleation requires higher

stresses at cryogenic temperatures than at room temperature.

For the triaxial stress state, an analysis similar to

that shown in Table 9 for the critical stress required to

nucleate voids is difficult. To develop a basis for discussion, look at the stress-strain history of a second-phase

particle outside of the process zone, PZ. The critical

location outside the PZ should be where the maximum normal

stress is found. McMeeking [24] found that the maximum

normal stress was at a distance of 2 or 3 times the value of

6. His results also show that the peak stress coincides

with a relatively low strain, approximately 0.01 (see Fig.

4) . For this discussion, the peak in maximum stress is

taken as the position ahead of the crack where the strain is

0.02 because 0.02 strain is the smallest strain that can be

estimated from the aDPH gradients. The stress-strain

197
Table 9. Data Used for Calculation of ac (MPa)

Required for Void Nucleation in 22-13-5.

Critical

Point

uniaxial

_JIeq O.Q. °c(l) Oc(2)

voids nucleate

from sigma at

295 K

1060 350 1060 425 1410 2075

voids nucleate

from microalloy

particles at

295 K

1820 1080 1630 425 2710 4065

voids nucleate

from sigma at

76 K

2860 1240 2810 1065 4050 5650

neck formation

due to local

heating at 4 K

2065 865 2025 1360 >2890 >3130

Biaxial

voids nucleate

from sigma

at 295 K

880 440 690 290 1130 1520

shear fracture

at 76 K

1990 990 1720 895 >2710 >3310

shear fracture 2080 1040 1800 1640 >2840 >2340

at 4 K

is interpreted in terms of a void-nucleation criteria and

called the interfacial stress, o\. The next two paragraphs

outline a method of estimating ay, o*q and aj for a distance

ahead of the crack where the strain is 0.02.

As the applied energy, J., increases, the size of the

deformed zone around the crack increases. The location

ahead of the crack tip where the strain is 0.02 can be

estimated from the results of the aDPH gradient. The hard-
ness gradient ahead of the crack tip has been measured at

each test temperature as a function of applied energy, Figs.

46, 53, and 61. The results of Fig. 38, where the strain in

unaxial and biaxial stress states was related to measured

aDPH, indicate that 0.02 strain is about 20 aDPH. For

example, in Fig. 46b for specimen AA-2, a smooth curve fit

through the aDPH data for 0=0 reaches 20 aDPH about 0.78

mm from the zero point, the tip of the original fatigue

precrack. The critical distance for AA-2 is 0.78 mm. The

next step is to draw a new modulus line, slope = 200 GPa, on

Fig. 17 at an offset of 0.02 strain. The value of a«q at

each temperature is given by the intersection of the new

modulus and the flow curves for the different temperatures.

Given the Ja at any point in the test, the maximum

stress, <7i, in the elastic region ahead of the crack tip can

be estimated from the results of the analysis of small-scale

yielding. The FEM results shown in Fig. 62 were used for

the calculation of a\. For any point inside the PZ, the

small-scale-yielding results are invalid. Given a value of

a*, and <7i, then the value of <7 t can then be calculated according to the equation,

<7 T = <71 - {1/ 3} Otq (8)

of Anderson [58]. The purpose of calculating the <7i is to

compare the local values of o\ in the toughness test to the

comparable values of oc for void nucleation calculated in

the biaxial and uniaxial stress states.

The results of the calculation of <7i at 4 K are shown

first because the 4-K tests show the least amount of plasticity. The critical distance ahead of the crack tip in the

interrupted C(T) specimens, A-10 and A-ll, is 0.18 and 0.10

respectively. A third value of Ja in the calculations is

taken to be equal to Ji c for specimen A-4 for which the

hardness gradient was taken; see Fig. 61c. The critical

distance r = 0.18 mm for A-4.

The data used in the procedure are shown in Table 10

and the results of these calculations in terms of criteria 1

and 2 are shown. The results for criterion 1 are plotted in

Fig. 63 as a function of Ja. The values of <jc calculated for

shear fracture in the uniaxial and biaxial stress states are

shown for comparison. The value of <7i must start at 0

because it is the sum of two parameters related to the

applied force. The value rises rapidly to 4.2 GPa,

199
r / (J/aa

o

€

o

)

Figure 62. Plot of the maximum normal stress ahead of

the crack tip in a strain hardening material

calculated as a function of applied J [54].

Table 10. Summary of

from

data for calculation of o\ at 4 K

results of toughness testing

Distance Ueq 0*1 Oj Crit. 1 Crit. 2

ahead of crack (MPa} (MPa} (MPa} (MPa}

e j. = 51 N/mm

r = 0.10 mm 1600 3540 2620 4220 3980

e j. = 94 N/mm

r = 0.18 mm 1600 3540 2620 4220 3980

2 J. = 59 N/mm

r = 0.18 mm 1600 3110 2180 3780 3550

assuming YS = 1325 MPa.

200
CL

O

cn

co

LU

DCH

CO

<

o

<

Liar

LU

h6

. , . T —— -i r

T = 4 K

5

A A

4

- — A

/

3 / Neck Formation • Uniaxial . — — — — — — —/— — — — = = = = = = = = = = = = :

¥ \\ Shear Fracture • Biaxial

2 /

/

/

1 /

- /

/

n / 1

, i i j 0 20 40 60 80 100

APPLIED ENERGY, N/mm

Figure 63. Plot of the interfacial stress (a i) , calculated according to Argon and Im [14], ahead of

the crack tip in compact specimens at 4 K of

22-13-5 as a function of applied J. Values

of <Jc for uniaxial and biaxial stress states

are shown for comparison.

decreases slightly to 3.8, and then increases back to 4.2

GPa. The maximum value of o\ at each of the three given Ja is greater than the value of ac from uniaxial and biaxial

stress states where no significant void nucleation occurred,

see Table 9. These results appear to be reasonable.

At 295 and 76 K, the load-displacement curves for the

complete tests are as shown in Fig. 3. Large-scale yielding

conditions exist at the point where Ji c is measured (refer to

Fig. 20 for an example of a test at 76 K) . Despite the

plasticity observed in the P-u curves, the deformed zones

ahead of the blunted crack tip are approximately 2 times <5

Calculations similar to those in Table 10 for the 4-K data

were performed on the comparable specimens (AA-4, AA-2, and

AA-1 at 295 K and AAA-5, AAA-4 , and AAA-2 at 76 K) and the

results are shown in Tables 11 and 12. Again, the a\ calculated by criterion 1 are plotted as a function of Ja in

Figs. 64 and 65.

201
Table 11. Summary of data for calculation of a

\

at

295 K from results of toughness testing.

Distance

ahead of

ahead

crack

Oeq

(MPa)

oy

(MPa)

at Crit

(MPa}

. 1 Crit. 2

(MPa}

£_J«_= 100 N/mm

r = 0.15 mm 520 2070 1770 2290 2220

£_J«_= 228 n/mm

r = 0.78 mm 520 1710 1410 1930 1860

= 524 N/mm

r = 1.38 mm 520 1820 1520 2040 1970

assuming YS = 425 MPa

Table 12 . Summary of data for calculation

76 K from results of toughness

of <7i at

testing.

Distance

ahead of

ahead

crack

a«q

(MPa}

o1 (MPa)

oj Crit

(MPa)

. 1 Crit. 2

(MPa)

§ j. = 54 N/mm

r = 0.21 mm 1200 2650 1960 3160 2870

§ j. = 187 N/mm

r = 0.21 mm 1200 3690 3000 4200 3910

§ J. = 330 N/mm

r = 0.31 mm 1200 3780 3090 4290 4000

assuming YS = 1066 MPa

202
INTERFACIAL

STRESS,

GPa

g

INTERFACIAL

STRESS,

GPa

4

3

2

1

0

T = 295 K

A

Void Nucieation - Nb, V ppt. Uniaxial

S'

/ a

*

/

• / Void Nucieation - sigma, Uniaxial

-/- - Void Nucieation - sigma, Biaxial

i

i

0 100 200 300 400 500 600

APPLIED ENERGY, N/mm

. 64 Plot of o, ahead of the crack tip in C(T)

specimens tested at 295 K.

APPLIED ENERGY, N/mm

Fig. 65 Plot of Oj ahead of the crack tip in C(T)

specimens tested at 76 K.

203
At 76 K, the values of o\ calculated from the toughness

tests increase steadily to 4.2 GPa at a Ja of 187 N/mm. The

value increases slightly to 4.3 GPa at JIc for specimen AAA2. The o i for the lowest Ja is below the value required to

nucleate voids from sigma in the uniaxial stress state. At

a slightly higher Ja , the a

i

is large enough to nucleate

voids from sigma phase. The value of oc for void nucleation

from microalloy particles is greater than 4.05 GPa.

At 295 K, the a i values at Ja are greater than the ac for void nucleation from sigma phase in uniaxial stress

state (1.4 GPa) . The maximum value of a is below ac for

void nucleation from microalloy particles (2.7 GPa).

The procedure for calculating o\ allows for useful

comparison of the different micromechanisms of fracture for

22-13-5. The values of a\ reach a plateau and do not increase significantly after a point. The plateau value is a

by-product of the procedure, which is limited by the smallscale-yielding analysis. The small-scale-yielding assumption does not hold after the ligament is yielding in the P-u

curve. A three-dimensional, large-scale-yielding analysis

of the peak stress would probably that show the value of o\

increases continually with Ja. The procedure described here does not relate to the

value of o\ between second-phase particles and the matrix

inside the PZ. The local value of ac for void nucleation at

microalloy particles appears to be met inside the PZ soon

after the specimen shows macroscopic yielding. Observations

of subcritical crack growth below Ji c in Figs. 42, 51, and 55

support this conclusion.

5.3 Discussion of models

Existing models attempt to characterize MVC in toughness testing by observations after the test is complete [30-

31]. The local strain to failure is assumed to be a function of the morphology of the fracture surface, SZ or dimple

shape. For 22-13-5, the formation of a SZ and dimples on

the fracture surface do not relate directly to toughness. A

crack can form at Ja < Ji» A key question is, Why does the

crack grow at a low rate, (dJ/da) , for Ja < Jic and at a

higher rate for Ja > Ji£

The measured toughness is directly related to a given

decrease in compliance of the specimen (a global parameter)

,

which is a function of the sensitivity of the experimental

204
procedure, and not necessarily to the local parameters at

the crack tip. The distribution of plastic strain around

the crack tip could explain why the formation of a crack

inside the PZ does not directly lead to detectable increase

in compliance. The best example to illustrate this point is

Fig. 9 in the introduction. AISI 304 austenitic stainless

steel tested at 76 K transforms to martensite when plastically strained. The etch has clearly defined the regions

of high strain because the martensite is attacked heavily.

With a different etch, a crack can be seen at the crack tip

for Ja = 0.33 Ju> When the specimen is unloaded to measure

its compliance, the plastically deformed regions around the

crack tip can exert a closure stress on the sharp crack.

The hardness profiles around the crack tips for 22-13-5 (see

Figs. 46, 53, and 61) show that a large plastic zone is

concentrated around the tip in a wake. This wake moves

along with the crack and could serve to close the crack when

the specimen is unloaded. This wake effect can be thought

of as similar to plasticity-induced crack closure in fatigue

crack growth experiments [60].

The analysis of a i for different micromechanisms leads

to a model for toughness based on nucleating voids in the

highly stressed, elastically constrained region outside the

PZ. The critical point is defined at the J« where the o\ for

void nucleation equals ac for void nucleation outside the

PZ.

If the value of a i reaches ac inside the PZ, voids can

form, grow, and coalesce to form a subcritical crack. The

plastic zone will exert a compressive stress on the crack

tip so that on unloading, the true crack length is not

recorded. If o\ * ac outside the PZ, then the voids can

nucleate. Once voids have nucleated, it is easy for the

voids to grow and coalesce with no additional J«. The

increment of crack growth is relatively large and measureable by the unloading compliance technique. Fracture toughness is then defined as the point where the oi (a function

of J« and the material's strain hardening characteristics),

equals the material's resistance, oc, at a point outside of

the PZ

The value of oc depends upon many factors and would be

difficult to calculate from first principles. Argon and Iro

[14] found that ac was a function of the local spacing of

void-nucleating particles and correlated with the material '

elastic modulus. In addition, the strain hardening characteristics should also be important in determining oe. The

205

win mi — —mm ——
value of oc is relatively independent of stress state so

that it can be experimentally measured in a uniaxial tensile

test and used as a failure criterion in the toughness test.

Observations from this test program support the idea

that void nucleation is the most important micromechanism

affecting toughness and that void growth is not directly

related to toughness. If void growth were important in the

determination of toughness, then the fracture surface appearance of the specimens at different test temperatures

should reflect the measured toughness. In fact, the fracture surfaces of toughness specimens tested in the T-L

orientation at different temperatures (Figs. 45, 52, 58, and

59) are remarkably similar despite the difference in toughness. With L-T test orientation at 4 K, the fracture surface appears more equiaxed, even though the toughness is

approximately the same as for the T-L specimens at 4 K.

The revised model for toughness described above is

based on the applied stresses for void nucleation from

microalloy particles. The initial model [44] was based on

the resistance of the material to void nucleation at widely

spaced inclusions. The definition of the critical point

that characterizes toughness in the revised model depends on

the applied stresses outside the PZ. In the initial model,

the critical point depended on the local stresses that exist

near an inclusion.

In order to predict the test results, a distance parameter is needed. The characteristic distance, L, is a key

feature of most models for toughness by either ductile or

brittle fracture mechanisms. Typically, the value of L for

ductile fracture is taken as the average spacing of voidnucleating particles. In brittle fracture, the value of L

is related to the value of 6 at which a suitable nucleation

site for cleavage enters the PZ [2]. The average inclusion

spacing, like the distribution of cleavage nucleation sites,

is related to the resistance of the material to void nucleation and does not change as J* increases.

If a steel is tested at a high test temperature, the

flow strength is low, the PZ is relatively large, and may

includes many inclusions. A correlation between inclusion

spacing and toughness may be appropriate. For the 22-13-5

steel at 295 K, the PZ extends about 0.6 mm ahead of the

crack tip. Since the average inclusion spacing is about 0.1

mm, the larger particles (inclusions) could play an important role in determining ac and toughness.

206
If the same steel is tested at a lover temperature, the

flow strength is higher, the PZ is smaller, and may not

include significant numbers of inclusions. In this case,

the average spacing of inclusions can not be fundamentally

related to toughness. For 22-13-5 steel at 4 K (HI-4K) , the

PZ is about 0.1 mm, or equal to the average inclusion spacing. Because there are few inclusions found in the PZ, it

is difficult to imagine the same relationship between average inclusion spacing and toughness that is found at higher

test temperatures.

The values of J« and 6 relate directly to the distance

ahead of the crack for which the stress and strain are

raised over the values in the bulk. The micromechanisms for

fracture can only operate within the PZ (2 to 36)

The

value of L is then defined in terms of the applied stresses

for void nucleation rather than in terms of the resistance

of the material. L could be related to the inclusion spacing for some test materials and limited test conditions, but

not in general. A similar conclusion was reached by others

[34].

The Of i and e* terms from equation 3 are not needed in

this analysis of void nucleation. The failure criterion is

defined only at the edge of the PZ where the plastic strain

is low. This fact simplifies the analysis, but technically

restricts the model to small-scale yielding or Kic testing.

A better analysis for the location of the peak stress would

improve the method. The limiting case would be if o\ never

reached oc outside the PZ. This would represent the situation where a different micromechanism interrupts the process, like in cleavage fracture. This frequently occurs in

toughness testing of ferritic steels [57,60], particularly

at high strain rates.

5.4 Material properties of 22-13-5

A key question in this work is, What effect do metallurgical variables like yield strength, second-phase-particle distribution, and alloy content have on ductile

fracture? The observations reported here form a basis to

evaluate these factors.

The best example is the role of the large, scattered

second-phase particles like sigma and nonmetallic inclusions. They are most important when the deformed zone is

largest, that is when the test temperature is highest and

207
the stress state is uniaxial. Then, a large volume of material is deformed and many particles are sampled. The

particles do change the stresses locally in the matrix (Fig.

39b) , and the location of final fracture is probably associated with the largest particles. However, the global effect

of large, scattered particles on the fracture stress is

minimal so the resistance of the steel to void nucleation is

not altered.

The results of the calculations of ai for the triaxial

stress state are consistent with the idea that the inclusions and sigma do not affect toughness. The value of o\

reaches ac for void nucleation from the sigma particles at

low J«. If void nucleation from sigma influenced toughness,

then Jic would be lower than the values measured here.

The calculation of at is not affected by the presence

of the large particles because o\ is a function of the yield

strength and macroscopic strain hardening, factors the

particles do not influence. In theory, a given distribution

could change the stress state locally from triaxial to

biaxial and reduce the constraint ahead of the crack tip.

This would lower the applied stresses.

Test temperature has a large effect on the yield

strength of 22-13-5. Decreasing the temperature from 295 to

4 K increased the yield strength by a factor of three. This

directly raises the applied stresses by an equal amount. A

similar effect could be ascribed to any strengthening mechanisms, like grain refinement or precipitation strengthening.

Any local variation in yield strength would cause a similar

local variation in o\. As the volume of material that is

strained decreases, the likelihood of yield strength variations increases. These two factors imply that larger

scatter in toughness can be expected at 4 K compared to test

results at higher temperatures.

Decreasing the test temperature from 295 to 76 K for

22-13-5 raised the resistance of the steel to void nucleation, ac$ in uniaxial stress state. At 4 K, local heating

and shear fracture interrupt the fracture process in uniaxial and biaxial stress states so that it is not clear if

oc at 4 K increases or not. The product of strength and

toughness, QI, does peak at 76 K so that qualitatively, the

mechanical performance also peaks at 76 K. It could be that

specimen heating at 4 K is affecting the material properties

rather than a direct consequence of the resistance of the

steel to void nucleation.

208
The microalloy particles do not affect the yield

strength of 22-13-5 directly. Indirectly, the presence of

microalloy particles influences the grain coarsening resis

tance of the steel and the interstitial content of the

matrix. Both factors relate to the yield strength and

therefore influence the driving force for void nucleation.

The substitutional alloying elements like Cr, Ni, and

Mn were constant in this research program, so there are no

direct observations on alloying effects can be discussed.

6.0 CONCLUSIONS

1. Both shear fracture and MVC are important micromechanisms of ductile fracture in 22-13-5. Where MVC is

observed, void nucleation absorbs considerable energy

and contributes significantly to the total energy

absorbed in ductile fracture. When shear fracture

occurs, it can be assumed that the energy required to

nucleate a significant volume fraction of voids is

greater than the energy required to cause shear fracture.

2. In fracture toughness testing, the first observation of

MVC on the fracture surface does not correspond to the

measured fracture toughness. The calculations of the

applied stresses related to void nucleation ahead of

the crack tip suggest that fracture toughness may be

described by the point where a critical stress for void

nucleation is reached at a characteristic distance

ahead of the crack tip instead of at the crack tip.

3. The characteristic distance related to toughness is a

function of the applied stresses for void nucleation in

this steel. This is in direct contrast to the generally accepted approach which is to relate the characteristic distance to the inclusion spacing, the resistance of the steel to fracture.

4. The applied stress for void nucleation in fracture

toughness testing (ai) is a strong function of the

yield strength, which for 22-13-5 depends mainly on

test temperature and interstitial nitrogen content.

Variations in nitrogen content will produce scatter in

the measured yield strength and fracture toughness for

different locations in the plate.

209
5. The critical stress, oc, for void nucleation can be

measured in a simple, uniaxial tensile test and can be

related to the critical value of fracture toughness (Ki c or

6.

The small microalloy particles in the austenitic matrix

influence the properties of the steel indirectly in two

ways. First, the particles help resist grain coarsening during thermal processing. A fine grain size

will increase the yield strength of the steel and its

resistance to strain localization. Second, the particles contain nitrogen, which decreases the nitrogen

content in solution and lowers the yield strength.

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3. J.W. Morris Jr. and E.N.C. Dalder, "Cryogenic

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216
WELDING

LEADER:

CONTRIBUTOR

T.A. Siewert

C . N . McCowan

WELDING PROGRAM

OBJECTIVES

o Investigation of the metallurgical factors that affect the mechanical

properties of stainless-steel weldments at cryogenic temperatures,

o Contribution to the development of improved filler metals for welding

stainless steels for use at 4 K.

o Evaluation of the mechanical properties of weldments at 4 K.

o Development of methods for detecting defects and evaluation of their

significance in stainless steel weldments.

RESEARCH HIGHLIGHTS:

o We evaluated a series of welds produced with a type 904L electrode

(20Cr-25Ni-4. 5Mo) . They had the highest toughness values which we have

measured for high strength welds (near 1000 MPa) at 4 K.

o A series of shielded metal arc compositions based on a type 308

stainless steel were unable to match the toughness of type 316LN and

21Cr-6Ni-9Mn base materials, even at 76 K.

o A study of a type 316L gas metal arc weld revealed a low energy

interdendritic fracture mode as nitrogen was added to reduce the

ferrite content and increase the strength.

219
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS

Weld Alloy

[designation]

Welding

Process

Tensile Charpy

Impact

Fracture

Toughness

Fatigue Crack

Growth Rate

* Report Volume

(page)

AUSTENITIC STAINLESS STEELS

Fe-13Cr-13Nl-20Mn-0 .2C SMAW XI XI XI (171)

Fe-13Cr-20Ni-9Mn-2Nb SMAW I. II I, II, I (159,169)

VIII II (299)

VIII (29)

Fe-15Cr-12Ni-10Mn SMAW IV IV, VIII IV (453)

VIII (29)

Fe-16Cr-15Ni-4Mo SMAW IV IV IV (275)

Fe-16Cr-16Ni-2Mo FCAW II III III, VI III (155)

VI (199)

Fe-16Cr-16Ni-4Mo FCAW IV IV, VIII IV IV (275)

VIII (29)

Fe-16Cr-35Ni-4Mn SMAW I I, VIII II I (159)

[330 modified] II (299)

VIII (29)

Fe-1 7Cr-9Ni-Mn-N SMAW IX IX IX (217)

Fe-17Cr-16Ni-4Mo SMAW IV IV IV (275)

Fe-18Cr-llNi-2Mo-N EBW X X X X (299)

LBW X X X

Fe-18Cr-l6Ni-6 . 5Mn-2Mo SMAW IX IX IX (257)

Fe-18Cr-16Ni-9Mn-N GMAW III, IV rv III, IV III, VI III (155)

IV (275,453)

V (199)

Fe-18Cr-18Ni-2Mo-5Mn-N SMAW IV IV IV (453)

Fe-18Cr-20Ni-5Mn-N GMAW IX, XI IX, XI IX (247)

XI (247)

Fe-18Cr-20Ni-6Mn-0 .3Nb GMAW rv IV IV IV (275 )

Fe-l8Cr-20Ni-6Mn-0 .2Ti GMAW IV IV IV IV (275)

Fe-l9Cr-12Ni-2Mo SMAW ii II, VIII II II (315)

[316] VIII (29)

Fe-19Cr-12Ni-2Mo EBW rv IV IV, VI IV (275,415)

[316L] VI (199)

GMAW III, IV, III, IV, IV, XII III (155,195)

XII VIII IV (303)

VIII (29)

XII (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); NBSIR 82 -1667 (V); NBSIR 83-1690 (VI);

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

NISTIR 89-3931 (XII)

220
SUMMARY Of MECHANICAL PROPERTY MEASUREMENTS OP WELDS, continued

Weld Alloy

[designation]

Welding

Process

Tensile Cherpy

Impact

Fracture

Toughness

Fatigue Crack

Growth Rate

Report Volume

(page)

Fe-19Cr-l2Ni-2Mo GTAW III, IV, III, IV. IV III (155,195)

[316L], continued XI VIII XI IV (303)

VIII (29)

XI (215)

SAW III, IV, III, IV, III, IV III, VI III (155,195

V V IV (275,415)

V (233)

VI (199)

SMAW II, III. II, VIII II, III. III, IV, VI II (299,315)

IV IV III (155,167)

VI (199)

VIII (29)

Fe-l9Cr-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-10Nl PCAW III III III, VI III (155)

[308L] VI (199)

GMAW III III III, VI III (155)

VI (199)

GTAW XI XI XI (215)

SMAW II, III II, VIII II, III III, VI II (299)

III (155)

VI (199)

« VIII (29)

Fe-20Cr-10Nl-N PCAW IV IV IV, VI IV (275,415)

[308LN] V (199)

Fe-20Cr-10Ni-9Mn-N SMAW I I I (159,169)

Fe-20Cr-l6Ni-7Mo-2W SAW IV IV IV (275)

Fe-20Cr-34Ni-2Mo-Nb SMAW VII VII VII (253)

Fe-21Cr-6Ni-9Mn-H GTAW XI XI XI (215)

SMAW I, II I. II I (159,169)

Fe- 2 5 C r-2 2N i-4 Mn- 2Mo GMAW XII XII

IIXII

(299)

( 187)

MATERIALS STUDIES POR MAGNETIC PUSION ENERGY APPLICTIONS AT LOW TEMPERATURES: NBSIR 78-884 (I);

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

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

NISTIR 89-3931 (XII)

221
SUMMARY OP MECHANICAL PROPERTY MEASUREMENTS OF VELDS, continued

Veld Alloy

[deslgnetlon]

Voiding

Process

Tensile Charpy

Impact

Fracture

Toughness

Fatigue Crack

Growth Rate

* Report Volume

(page)

AUSTENITIC STEELS

Fe-15Mn-8Ni-lMo-0 . 7C SAV VIII VIII VIII (15)

Fe-5Cr-25Mn-lNi V V.VII V.VI V (233)

VI (199)

VII (245)

Fe-6 . 5Cr-25Mn-3Ni-lMo SMAV V V V.VI V (233,245)

VI (199)

ALUMINUM ALLOYS

Al-5Mg

[5183]

GMAV III IV IV III (155,217)

IV (323)

Al-5Mg-0.12Ti

[5556]

GMAV III IV IV III (155, 217)

IV (323)

Copper Alloys

Cu-0.08Ag

[C10700]

XI XI (201)

Cu-0.4Be-l.8Ni

[C17510]

XI XI (201)

MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICTIONS AT LOV TEMPERATURES: NBSIR 78-884 (I);

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

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

222
DEVELOPMENT OF AN SMA ELECTRODE

TO MATCH TYPE 316LN BASE METAL CRYOGENIC PROPERTIES

T.A. Siewert and C.N. McCowan

Materials Reliability Division

National Institute of Standards and Technology

Boulder, Colorado

A series of shielded metal arc electrodes was fabricated and the

mechanical properties of the resultant welds were measured. The

compositions are based on a high-Mn, high-N version of a type 308 electrode,

to which additional Mn, N, Ni, and Mo were added. The strength was measured

at 4, 76, and 298 K, and the Charpy V-notch absorbed energy was measured at

76 K. The strengths of the welds met or exceeded the values for the base

metal, but the Charpy V-notch absorbed energies fell short.

BACKGROUND

The trend toward more efficient designs for structures continues, with

designers seeking more strength and toughness from the structural materials.

This general observation is true even for the special category of cryogenic

structures. The Compact Ignition Tokamak being designed by Princeton Plasma

Physics Laboratory requires property data on such high strength materials

for the service range of 76 K to 298 K. Designs for other cryogenic

structures are also expected to require a combination of high strength and

toughness, to temperatures as low as 4 K when used with conventional

superconducting magnets. Thus, it is useful to collect information on high

strength candidate materials, in an attempt to understand how the properties

can be improved further.

Whereas several stainless steel plate materials, such as 21Cr-6Ni-9Mn

and type 316LN, demonstrate properties that might be suitable for cryogenic

structures, the welding electrodes necessary to join these compositions into

a fabricated structure remain poorly characterized. Early studies developed

screening data on high strength weld compositions, but the volume of such

data is limited.

Espy compared the properties of the matching composition electrode for

21Cr-6Ni-9Mn with various experimental compositions, at both room

temperature and cryogenic temperatures. 1 Although the yield strength

increases sharply as the temperature drops, his data in Table 1 also show

that the toughness of these welds at liquid nitrogen temperature is marginal

when compared to a common design requirement, such as 20 J of absorbed

energy in Charpy V-notch (CVN) testing. Such a toughness requirement

becomes more important as the thickness of the structure increases.

223
Table 1. Mechanical Properties of Variations on a 21Cr-6Ni-9Mn Weld.

Mechanical Properties Mechanical Properties

at Room Temperature at 76 K

Yield CVN absorbed Yield CVN Absorbed

Composition Strength (MPa) Energv (J) Strength (MPa) Energv (J)

21Cr-6Ni-9Mn-0 . 2N 570 54 1120 14

13Cr-20Ni-9Mn-0. 18N 525 36 820 16

20Cr-10Ni-9Mn-0. 19N 560 66 945 13

18Cr- 16Ni-9Mn-0 . 15N - - 880 24

Espy attributed the poor toughness shown in Table 1 to the high

strength (due to nitrogen additions). This causes a problem for structural

designers since the high strength is precisely what makes the 21Cr-6Ni-9Mn

composition attractive, and produces benefits such as reducing the required

section thickness. The Espy study also contained data for welds having the

same combinations of Cr, Ni, and Mn but with lower concentrations of N.

These compositions had better toughness (up to 38 J) at 76 K, but the

strength dropped (as low as 640 MPa)

.

The 21Cr-6Ni-9Mn base metal has much better properties than the welds

(similar strength to the 21Cr-6Ni-9Mn weld metal at the two temperatures

with CVN absorbed energies of 280 J at room temperature and 88 J at 76 K)

.

This good combination of strength and toughness in the base metal could lead

a designer to make incorrect assumptions about the weld properties. Unless

the structural designers are aware of the lower toughness of the welds, they

might assume similar properties for the welds and base metal, and jeopardize

the integrity of the structure. Perhaps the most important unanswered

question is: if the various weld compositions have toughness values that are

inferior to the base metal, how can a high toughness structure be

constructed?

Other austenitic stainless steels and matching weld compositions, such

as AISI type 304 (19Cr-9Ni) and type 316 (18Cr-12Ni-2Mo) , have received a

more thorough evaluation at cryogenic temperatures. 3

’ 4 One study of the

welds (Ref. 2) shows a linear increase in strength (measured at 4 K) as the

nitrogen is increased from 0.04 to 0.28 wt.%. The study presented the CVN

absorbed energy for the same alloys at 76 K and showed a reduction in the

absorbed energy (one measure of toughness) from 41 to 20 J as the strength

increased. These data indicate that the strength and toughness are

inversely related for these austenitic stainless steels, as they were for

the 21Cr-6Ni-9Mn. Another study (Ref. 3) reports such an inverse

relationship between strength and toughness in both type 304 base metal and

types 308L (20Cr-10Ni) and 316L weld metal, as the alloys are alloyed with

N. The 4 K yield strength versus 4 K fracture toughness is shown in Figure

1. This relationship implies that a designer can modify the strength and

the toughness by varying the nitrogen content, but that a choice of higher

strength implies lower toughness and vice versa.

224
Figure 1. Fracture toughness versus yield strength for velds and base metal

at 4 K.

The strength- toughness trend has been used to define a higher strength

stainless steel base metal, type 316LN. This alloy has a similar

composition to type 316 (18Cr-12Ni-2Mo) , but has the carbon restricted below

0.03 wt.% and a nitrogen range of 0.10 to 0.16 vt.%. This nitrogen level

strengthens this alloy over that of type 316 or 316L, and reduces the

toughness. Hovever, the retained toughness is quite adequate for the

intended applications. The trend line for type 304 base metal in Figure 1 is also representative of type 316LN. For a strength of 1000 MPa, the trend

line indicates a toughness near 200 MPa*m"*/2 at 4 K. Notice that the velds

that match this alloy in strength have much lover toughnesses. When

evaluated at 76 K, the type 316LN base metal strength drops to near 730 MPa,

but the toughness increases, perhaps as much as 50 % over the 4 K value.

Therefore type 316LN base metal is also a good candidate for use at liquid

nitrogen temperature. Again, the most important question is, vhat electrode

can be used to match both the strength and the toughness of this base metal

at cryogenic temperatures, whether at 4 or 76 K?

225
One solution for applications requiring high toughness and strength is

presented in Ref. 4. This study shows that the 76 K CVN absorbed energy is

influenced strongly by the Ferrite Number (FN) and nickel content. The

toughness increases with nickel and decreases with FN. Since increasing the

nickel content is also known to reduce the FN, nickel additions improve the

toughness by both means

.

An alloy designer can use this information to construct new alloys. Of

course other factors must also be considered, but in simplest terms the

nitrogen level can be adjusted to achieve the desired strength and the

nickel level can be adjusted to change the toughness. Alloys with desirable

balances of properties do not exist yet because the market for welds in

thick cryogenic structures is still in its infancy and these compositionproperty relationships have been determined in detail only recently. The

program reported here was designed to develop electrodes having properties

that match type 316LN, or perhaps 21Cr-6Ni-9Mn.

PROCEDURE

A series of eight shielded metal arc welding electrodes were produced

by an outside contractor. Using each electrode, welds were produced in

accordance with AWS Specification A5.4. This specification also describes

the production of standard weld pads, for compositional determination. From

each weld, five CVN specimens (all tested at 76 K) and three 6-mm-diameter

tensile specimens were prepared (one for each temperature 4, 76, and 298 K)

.

The eight electrodes comprised a composition matrix with a base

composition near type 308. The matrix included significant variations in

the contents of Mn, Mo, N, and Ni to determine the optimum level for these

elements. In other words, should we chose a standard weld composition that

nearly matched the type 316LN base plate (near 18Cr- 10Ni-2Mo-0 . 14N) , or

would a significantly different composition provide some substantial

property benefits? The matrix consisted of two levels of molybdenum (0 and

3 wt.% to compare a type 308 to the upper limit of a type 316 composition),

two levels of nickel (to measure how much this addition increased the

toughness), two levels of nitrogen (0.16 and 0.24 wt.% to evaluate the

toughness at two strength levels), and two levels of manganese (3 and 6 wt.%

for the low nitrogen level and 6 and 9 wt.% to provide a higher solubility

level for the higher nitrogen level)

.

The base nitrogen level of 0.16 wt.% was chosen to match the strength

of the type 316LN base metal. Since little data is available at these

temperatures and for higher nitrogen alloys, the other elements were varied

to evaluate the effects of their interactions on toughness. Our composition

range is shown in Table 2. The identification system is based on which

elements in the matrix are at higher levels for a given composition.

226
Table 2. Shielded Metal Arc Composition Design Matrix.

Weld Content in Weight Percent

Identification Mn Bi Mo B

Ni 3 14 . 0.16

NiMo 3 14 3 0.16

Mn 6 9 - 0.16

MnMo 6 9 3 0.16

NiN 6 14 - 0.24

NiNMo 6 14 3 0.24

MnN 9 9 - 0.24

MnNMo 9 9 3 0.24

RESULTS and DISCUSSION

Table 3 lists the actual compositions from the veld pads, including

both the major elements and the trace elements.

Table 3. Weld Compositions.

Weld

ID

Content in Weight Percent

Mn

Ni 0.027 2.76 0.29 0.018 0.008 17.94 14.15 0.07 0.11 0.147 0.070 0

NiMo 0.029 2.67 0.31 0.019 0.006 18.06 13.66 2.66 0.09 0.147 0.066 0

Mn 0.034 5.82 0.31 0.019 0.009 19.58 9.17 0.07 0.11 0.177 0.052 1.5

MnMo 0.034 5.52 0.30 0.021 0.011 16.80 9.21 2.95 0.09 0.162 0.054 2.4

NiN 0.027 5.58 0.32 0.017 0.010 18.33 14.24 0.07 0.10 0.216 0.054 0

NiNMo 0.028 5.08 0.34 0.019 0.009 18.36 13.90 2.57 0.09 0.206 0.054 0

MnN 0.033 8.51 0.27 0.018 0.012 20.59 9.29 0.07 0.10 0.241 0.063 1.6

MnNMo 0.030 8.79 0.33 0.021 0.016 17.94 9.29 3.03 0.09 0.221 0.053 2.4

In most cases, the actual compositions in Table 3 are very close to the

desired values in Table 2. An exception was that the nitrogen content was

limited by reaching saturation in several of the veld pools. We concluded

that saturation had been reached because reformulation of the electrode

coating with a higher nitrogen level (to bring the nitrogen content closer

to the desired value) had no effect, other than to produce microporosity in

the weld pool. This microporosity did not appear on the tensile or CVN

fracture surfaces and so is not thought to have had a significant effect on

the properties. Notice that the chromium level fluctuates more among the

four welds with the lower nitrogen contents . This is because these alloys

were formulated to a ferrite value near 2 FN, which reduces the cracking

sensitivity for these alloys in the FA solidification mode (primary

austenite, then ferrite also).

227
comfortable margin over the strength of type 316LN base metal (76 K yield

strength near 735 MPa), showing that we reached this goal. Unfortunately,

the CVN values are far below the 88 J absorbed energy value for the 21Cr6Ni-9Mn base metal. (The absorbed energy value for the type 316LN base

metal is expected to be even higher.) The welds with the higher nickel

level are tougher, on the average, than those with the lower nickel levels,

but none of the compositions have a toughness high enough above the desired

lower bound value of 20 J to allow for the inevitable lot- to- lot and

procedural fluctuations that occur in the fabrication of large and complex

structures. The inferior toughness compared to what was desired is

attributed both to not increasing the nickel above the 14 wt.% level, and to

the higher inclusion contents inherent in welds produced with the shielded

metal arc process. In retrospect, we might have been able to fit several

more percent nickel into the coating formulation of the higher nickel

compositions, but even that may have been insufficient since the required

range may be as high as 20 to 25 wt.% Ni.

Table 6. Mechanical Properties of Shielded Metal Arc Welds at 4 K.

The 4 K data confirm the trends seen at 298 and 76 K. The strengths

continue to increase as the temperature is lowered to 4 K, and the specific

alloy compositions that were strongest at 298 and 76 K tend to be the

strongest at 4 K also.

Figure 2 summarizes the strength data for the three temperatures. This

figure clearly shows the increasing effect of nitrogen strengthening at

lower temperatures. This figure is also useful in identifying the nitrogen

level that is necessary to achieve a specified strength at a given

temperature

.

Yield Tensile

Weld Strength Strength Elongation

_ID_ (MPa) ..(MPa ) LU .

Ni 980

NiMo 1070

Mn 1040

MnMo 1160

NiN 1200

NiNMo 1180

MnN 1275

MnNMo 1295

1200

1400

1470

1390

1410

1315

1540

1455

12

30

31

11

12

5

17

2

228
0.10 0.20 0.30

Nitrogen Content, wt %

Figure 2. Yield Strengths versus Nitrogen Contents at

Three Temperatures

Figure 3 compares the 76 K strength and the 76 K absorbed energy. All

the data fall within the same trend band. Although the higher nickel

content is along the upper side of the band, this is much less of an

improvement in toughness that what we had expected from previous studies. 4 Therefore, it is difficult to predict what nickel level is necessary for the

desired absorbed energy level, but it is expected to be near the levels in

Ref. 4 (25 wt.% Ni)

.

The trend band has the expected shape. The highest strength

compositions have the lowest toughness.

229
ft-lb (-320 F)

LL

o

CM

CO

‘cn

Figure 3. Yield Strengths versus Charpy V-notch Absorbed Energy,

both at 76 K.

CONCLUSIONS

1. The shielded metal arc welds, formulated to match the properties of the

type 21Cr-6Ni-9Mn and 316LN base metals, matched their strengths, but fell

short of matching their toughnesses.

2. To match both the strength and toughness of the base metal, a

substantially higher nickel content is necessary. This demonstrates that

the standard rule of choosing a welding electrode composition that nearly

matches the base plate is not appropriate for cryogenic applications of

stainless steels with stringent mechanical property requirements.

ACKNOWLEDGMENTS

Support for the 298 and 76 K part of the program was furnished by

Princeton Plasma Physics Laboratory. Support for the 4 K part of the

program was furnished by the Department of Energy. The electrodes were

formulated and welded by Alloy Rods.

REFERENCES

1. R.H. Espy, Weldability of 21-6-9 Stainless Steel, Materials Studies

for Magnetic Fusion Energy Applications at Low Temperatures - I,

NBSIR 78-884, National Institute for Standards and Technology,

Boulder, Colorado, 1978, p. 169.

230
2. 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, 66, 1987, p. 84-s.

3. R.L. Tobler, T.A. Siewert, and H.I. McHenry, Strength -Toughness

Relationship for Austenitic Stainless Steel Welds at 4 K,

Cryogenics, 26, July 1986, p. 392.

4. T.A. Siewert, Predicting the Toughness of SMA Austenitic Stainless

Welds at 77 K, Welding Journal, 65, 1986, p. 23.

5. C.N. McCowan and T.A. Siewert, The Effect of Inclusions on the

Fracture Toughness of 316L Stainless Steel Welds at 4 K, Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures

- XII, NISTIR 89-3931, 1989, p. 199.

6. C.N. McCowan and T.A. Siewert, The Fracture Toughness of 25Cr22Ni-4Mn-2Mo Stainless Steel Welds at 4 K, submitted to Advances in

Cryogenic Engineering, Vol. 36, 1989.

231
-
CRYOGENIC MECHANICAL PROPERTY DATA FOR

20Cr-25Ni-4. 5Mo GAS METAL ARC WELDS

T.A. Siewert and C.N. McCovan

Materials Reliability Division

National Institute of Standards and Technology

Boulder, Colorado

A 20Cr-25Ni-4 . 5Mo gas metal arc electrode was welded with Ar-N

shielding gas mixtures to evaluate their effect on the cryogenic mechanical

properties. The welds were found to follow the expected trend of increasing

strength with increasing nitrogen content at temperatures between 4 and

298 K. At 76 K and the solubility limit of N in the weld (0.18 wt.%), the

welds were found to have the best combination of strength (830 MPa) and

Charpy V-notch absorbed energy (104 J) which we have ever measured for a

weld.

BACKGROUND

Welds with properties that match those of the base plate are necessary

if designers are to have maximum latitude in planning the fabrication of a

structure. This general statement is also quite applicable to structures

for high-strength cryogenic applications. Various plate materials have been

developed and evaluated in some detail for high-strength cryogenic

structures, but few welding electrodes have been developed to join these

plates into a fabricated structure. Also, the existing stainless steel

electrodes (primarily designed for high- temperature applications) have been

poorly characterized for cryogenic service.

Recent studies reveal some of the relationships that determine the

strength and toughness for these stainless steels when evaluated for

cryogenic service. One study [1] shows a linear increase in strength

(measured at 4 K) as the nitrogen content of type 308 and 316 welds is

increased from 0.04 to 0.28 wt.%. The study determined the Charpy V-notch

(CVN) absorbed energy for the same welds at 76 K and found a reduction in

absorbed energy, from 40 to 20 J, as the 4 K strength increased from 400 to

1200 MPa. Of course only the low-nitrogen end of the range contains

compositions that satisfy the nitrogen limits for these grades, but these

data do show (1) how to increase the strength and (2) that the strength and

toughness are inversely related for these austenitic stainless steels.

Another study [2] reports such an inverse relationship for 4-K strength and

4-K toughness (fracture toughness) in both type 304 base metal and types

308L and 316L weld metal. These studies both imply that a designer can

choose either higher strength or higher toughness by varying the nitrogen

content, but that a choice of higher strength implies lower toughness and

vice versa.

233
Advanced designs for superconducting coils supports require highstrength materials. Some early studies developed screening data

specifically for high-strength veld compositions, but the volume of such

data is very limited. Espy [3] compared the properties of the matching

composition electrode for 21Cr-6Ni-9Mn-0.2N with various experimental

compositions, at cryogenic temperatures. He found that the toughness values

at liquid nitrogen temperature for these high-strength velds are marginal

when compared to a common design requirement, such as 20 J of absorbed

energy in Charpy V-notch (CVN) testing. Espy attributed the poor toughness

to the high strength of these electrodes (due to nitrogen additions).

This causes a problem for structural designers since the high strength

that justifies a reduction of wall thickness is precisely vhat makes the

21Cr-6Ni-9Mn composition particularly attractive for cryogenic applications.

The 21Cr-6Ni-9Mn base metal (unwelded) has much better toughness (CVN

absorbed energies of 88 J at 76 K) than the velds of similar strength.

One solution for applications requiring velds with both high strength

and high toughness is presented in Ref. 4. This study shows that the 76 K

CVN absorbed energy is determined primarily by the Ferrite Number and the

nickel content. Since increasing the nickel content is known to also reduce

the Ferrite Number, nickel additions can be used to improve the toughness.

The present study evaluates such a high -nickel composition, when

alloyed with nitrogen to produce high strength.

PROCEDURE

Since we did not have stifficient resources or time to melt a series of

alloys in a suitable test matrix, we used a welding technique to modify the

nitrogen content of a GMA electrode with a composition near the desired

value

.

We obtained a gas metal arc welding (GMAW) electrode with a high nickel

content and low levels of the impurity elements that might promote hot

cracking. The electrode had a 1.2 mm diameter, with an American Welding

Society designation of E904L and a nominal composition of 20Cr-25Ni-4. 5Mol.SMn. The strength of this composition was increased by adding N to the

veld pool through the shielding gas. We produced a reference weld with a

pure argon gaa shield and then added nitrogen at Increasing levels until

porosity occurred in the veld. Such an approach is not recommended in

production because arc length changes affect the nitrogen recovery, but the

approach does provide a rapid method of surveying the mechanical properties

over a range in strength. We chose gas mixtures with 10, 15, and 20%

nitrogen by weight. Previous research has shown that these shielding gas

mixtures cause a large Increase in the strength, as N content increases to

the solubility limit of the weld pool [5]. This procedure had some

limitations, such as not allowing us to add Mn to Increase the N solubility,

but did allow us to develop screening data for a veld of the modified

composition.

We produced welds using a 60 degree included angle single -V- groove

joint in 25 -mm- thick stainless steel plate. We made the velds in the flat

234
position using 220 A and 30 V, current and voltage levels that produced the

best bead shape and droplet transfer. After radiography, ve machined

6 -mm -diameter tensile specimens and 10 -mm- square CVN specimens, both with

their test areas centered in the veld metal. The tensile specimens had a

25 -mm gage length, so the elongation values were computed on this basis.

The room temperature and cryogenic tensile specimens were evaluated by the

standard procedure of ASTM E 6; however the cryogenic specimens required a

special load frame to keep them at the test temperature [3]. The CVN

specimens were immersed in a liquid nitrogen bath (76 K) , then tested

according to the requirements of ASTM E 23.

RESULTS and DISCUSSION

Veld Nitrogen Contents

Table 2 lists the amount of nitrogen that was measured in the velds, as

a function of the shielding gas mixture. The poor correlation betveen the

mixture ratio and the amount of nitrogen reinforces the statement that this

is not a suitable technique for production velds. The irregularity is

attributed to variations in the arc length from one veld test plate to

another, which changes the amount of nitrogen dissolved in the veld.

However, for each veld plate, ve noted less than a 5% variation in the two

or three individual measurements that were averaged to give these values.

This consistency within a weld indicates that the values are representative

of each veld. The veld vith 0.27 vt.% nitrogen had some pores, indicating

that this veld exceeded the solubility limit for nitrogen.

Table 2. Nitrogen Contents (vt.%) for the Shielding Gas Mixtures.

Mechanical Properties

Table 3 lists the mechanical properties for the gas metal arc

electrode. The strength values compare favorably vith those reported by

Espy, but the toughness values are substantially higher. Therefore, the

plan to improve the toughness by higher nickel contents was successful.

The measured nitrogen contents for each veld are Included for

comparison. Notice the lack of an increase in strength as nitrogen is

increased above 0.18 wt.%. This indicates that nitrogen above this level is

not distributed in the matrix in a form that promotes strengthening, but is

probably in the form of pores. The data in Ref. 1 suggest that higher

Shielding

&£

Nitrogen

Content

pure Ar

Ar/10% N

Ar/15% N

Ar/20% N

0.06

0.16

0.27

0.18

235
strength would have been achieved If the composition had a higher solubility

for nitrogen (such as an electrode formulated with a higher manganese

level) . The optimum manganese level would be the level sufficient to

dissolve the nitrogen content that provides the desired strength.

Table 3. Mechanical Properties of 904L (20Cr-25Ni-4.5Mo-l.5Cu).

Temperature

no

Nitrogen

Content

Yield

Strength

(MPa)

Tensile

Strength

(MPa)

Elongation

(%)

Reduction

of Area

(%)

CVN

Energy*

(J)

298 0.16 385 640 40 63

m 0.18 420 675 47 70 .

ft 0.27 415 665 48 67 -

76 0.16 750 1160 55 48 139

19 0.18 830 1240 57 36 104

n 0.27 780 1170 50 48 147

4 0.06 800 1110 35 m m 0.16 940 1310 52 35 .

* Average of results from two or three specimens.

These CVN absorbed energy values are the highest we have ever measured

for a weld. (The slightly lower values at 76 K for the 0.27 wt.% N weld are

attributed to the porosity that was revealed on the fracture surface.) These

excellent values are attributed to the combination of a very tough and stable

austenitic structure, and the use of a shielding gas that is very low in

oxygen content. Our recent research has shown a strong correlation between

high toughness and low oxide inclusion contents [5]. Since most austenitic

stainless steel GMAV electrodes are welded with an Ar/2% 0 shielding gas, it

is reasonable that replacement of the oxygen with nitrogen will reduce the

veld oxygen content, and so the Inclusion content. An Ar/N shielding gas

mixture produces a, transfer mode and veld puddle fluidity different and less

desirable than an Ar/0 mixture, but the importance of good properties

dominates for this application.

Micros truetural Evaluation

In addition to measuring the mechanical properties, we investigated the

microstructure of the 904L velds. Figure 1 shows the structure of a tensile

specimen at a 20X magnification for the veld produced with the Ar/10% N mix.

The microstructure has the large grain size typical of a fully austenitic

structure, each grain containing many dendrites of the same orientation

(dendrite packets). The interdendrltic regions and grain boundaries are free

of the large second phase regions and precipitates seen in fully austenitic

velds of lower toughness [6]. The inclusion content, distribution, and

diameters are typical of austenitic velds produced with a low-oxygen welding

procedure

.

236
Figure 1. Microstructure of weld produced using a Ar/10% N

shielding gas. (20X)

Figure 2. Fracture surface of a tensile specimen from a veld

produced using an Ar/15% N shielding gas. (300X)

Fracture Surface Appearance

The fracture surfaces of the tensile specimens had macroscopic fracture

features associated with high toughness, and could be best described as

consisting of ductile dimples and shear zones. There was no evidence of

237
I

brittle fracture, and only several indications of fracture along the dendrite

cell boundaries (also known as ribbed or corduroy fracture). This behavior is

quite consistent with other welds that exhibited good toughness. Previous

studies of compositions with poorer toughness exhibited very large percentages

of brittle fracture or rupture along austenite grain boundaries. Figure 2

shows the ductile dimple appearance of the fracture surface at 300X, in this

case the fracture surface of a specimen tested at 4 K. This specimen shows

several examples of the corduroy fracture.

CONCLUSION

(1) This 904L weld exhibited excellent CVN absorbed energy values, the

highest values which we have ever developed for welds at 76 K.

(2) The combination of good toughness and strength make 904L a good

candidate material for cryogenic structures.

ACKNOWLEDGMENTS

This program was supported by Princeton Plasma Physics Laboratory. The

electrode was furnished by the Wire Division of Sandvik Steel.

REFERENCES

1. 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, 66, 1987, p. 84-s.

2. R.L. Tobler, T.A. Siewert, and H.I. McHenry, Strength 'Toughness

Relationship for Austenitic Stainless Steel Welds at 4 K, Cryogenics,

26, July 1986, p. 392.

3.

R.H. Espy, Weldability of 21 -6-9 Stainless Steel, Materials Studies

for Magnetic Fusion Energy Applications at Low Temperatures - I,

NBSIR 78-884, National Institute for Standards and Technology,

Boulder, Colorado, 1978, p. 169.

4. T.A. Siewert, Predicting the Toughness of SMA Austenitic Stainless

Welds at 77 K, Welding Journal, 65, 1986, p. 23.

5. C.N. McCowan and T.A. Siewert, The Effect of Inclusions on the

Fracture Toughness of 316L Stainless Steel Welds at 4 K, Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures - XII, NISTIR 89-3931, 1989, p. 199.

6.

C.N. McCowan and T.A. Siewert, The Fracture Toughness of 25Cr-22Ni4Mn-2Mo Stainless Steel Welds at 4 K, submitted to Advances in

Cryogenic Engineering, Vol. 36, 1989.

238
FRACTURE CHARACTERIZATION OF 316LN WELDS WITH INTERCELLULAR FERRITE at 4 K

C.N. McCowan and T.A. Siewert

Materials Reliability Division

National Institute of Standards and Technology

Boulder, CO 80303

The cryogenic strength and toughness of type 316LN stainless

steel welds were evaluated. The weld composition was produced by

adding nitrogen to an argon based shielding gas during welding

with a type 316L stainless steel gas metal arc welding (GMAW)

electrode. The solidification mode of the welds was primary

austenitic: intercellular ferrite contents are approximately 4

percent. The average yield strength (950 MPa), ultimate strength

(1400 MPa) and fracture toughness (135 MPa-m*5

) of the welds at 4 K

are typical of values found for austenitic stainless steel welds

at cryogenic temperatures. This report focuses principally on the

fracture paths observed in tensile and compact tension specimens,

and the role of ferrite in the fracture. The ferrite phase in

these welds deformed in a ductle manner during failure.

INTRODUCTION

The good strength and toughness of type 300 series stainless steels at

4 K makes them attractive for use as structural materials in superconducting

magnet designs using Nb 3 Sn superconductors [1]. In particular, type 316LN

plate material is currently considered a prime candidate for the fabrication

of superconducting magnet cases, and welds of matching composition to join

this plate material are an obvious choice [2]. Having identical

compositions for the weld and plate materials, however, does not necessarily

result in identical mechanical, physical, and corrosion properties.

The grain morphology, compositional homogeneity, and micros tructural

phase constituents in the welds and plate material differ significantly due

to differences in their thermal and deformation histories. Type 316 plate

microstructures are fully austenitic, typically have equiaxed grain sizes of

around 50 pm, and have a homogeneous composition [3-4]. In type 316LN

welds, the microstructures are often two-phase (austenite + ferrite) [5].

Also, differences between primary ferritic and primary austenitic

solidification modes affect the morphologies and location of the respective

phases in the microstructure, and modify the solute segregation [6-7].

Finally, the grain size and inclusion content of the welds are much greater

than those typical of stainless steel plate materials.

Past studies show type 300 stainless steel welds have a broad range of

strength and toughness at 4 K [8]. In Fig. 1, the range in these properties

239
is indicated by a scatter band that is based on data from type 308L and 316L

welds. The individual data points identified in the figure are from

stainless steel GMA welds made to evaluate the effect of inclusion content

on toughness. For these welds, the inclusion contents varied over a range

similar to the inclusion contents of the welds used to construct the scatter

band (55 000 to 10 000 inclusions per mm2

) and they show that much of the

scatter in properties for the 308L and 316L welds is attributable to

differences in the inclusion contents [9]. The data for the fully

austenitic welds ( 18Cr- 20Ni- 5Mn-0 . 15N) in Fig. 1 also show the increased

toughness, due to the higher Ni content, and the increased strength, due to

the higher N content of these welds compared with the type 316L stainless

steel weld composition [10].

This study evaluates the mechanical properties of a low- ferrite -content

316LN weld composition that solidified in a primary austenitic mode. Some

ferrite was desired in the welds to increase their resistance to

solidification cracking [7]. The welds in this study were expected to be

tougher than 316L welds previously tested (Fig. 1) due to slightly higher Ni

content, lower ferrite content, and austenitic grain morphology. Increasing

inclusion content was expected to decrease the toughness of the welds.

Yield strength, MPa

Figure 1: The trend of decreasing toughness with increasing strength at 4 K

is indicated by the scatter band (dotted lines) developed using type 308L

and 316L stainless weld data (data not shown) [8]. The data points included

in the figure are from: (1) fully austenitic (18Cr-20Ni-5Mn-3Mo-0 . 15N) GMAW

welds (dots) in which inclusion contents range from 9 000 to 23 000

inclusions per mm2

, and (2) 316L stainless steel GMAW welds (18Cr-llNi-2Mn3Mo, FN-7) that solidified in an FA mode and have inclusion contents ranging

from 19 000 to 55 000 inclusions per mm2 [9].

240
MATERIALS AND PROCEDURE

The composition of the 316L GMA electrode used in this study is given

in Table 1. This composition was chosen because it has a sufficiently high

ferrite content (predicted to be 8 FN (approximately 8% ferrite) by the WRC

diagram [11], so that N could be added during welding, to increase the

strength of the weld, and some ferrite would still be present in the

microstructure. Ideally, we wanted to produce welds that solidified in a

primary austenitic mode, had ferrite contents near 4 or 5 FN, and had yield

strengths of 1000 MPa at 4 K.

Table 1. The GMA electrode composition.

c Mn P S Si Cr Ni Mo Co Cu Nb Ti

.014 1.63 0.014 0.010 0.38 18.6 12.2 2.58 0.03 0.08 0.01 0.01

To make the welds, a type 304L stainless steel base material was

beveled to provide a 60° single V- groove weld configuration. The plates

were approximately 300 mm long and 20 mm thick: 12 weld passes were needed

to fill the joint. Welding was done at 27 V and 230 A with an electrode

extension of 16 mm. A shielding gas of 10% N and 90% Ar was used. The

weld identified as SN was made using a conventional GMA shielding nozzle and

a Ar-10%N shielding gas. The DN weld was made using double Ar-N shielding:

a second, larger shielding nozzle was fitted around the usual shielding

nozzle to provide improved shielding of the arc. This reduces the amount of

oxidation that occurs during welding and thereby lowers the inclusion

content of austenitic stainless steel GMA welds [9],

From each weld, one all-weld-metal tensile specimen and two compact

specimens, notched in the direction of welding, were machined for testing at

4 K. Testing procedures and equipment are described elsewhere [12].

One compact tension specimen from each weld was postcracked and

fractured following the determination of the Kjc (J) values, to examine their

fracture surfaces. The second specimen from each set was left intact

following testing, so that metallographic cross sections of the crack tips

could be ground and polished for evaluation.

RESULTS

Microstructure

The microstructure of the welds had cellular-dendritic features

indicative of primary austenitic solidification. Eutectic ferrite was

241
present at intercellular regions in the structure: magnetic measurements

showed ferrite content to be approximately 4 FN. The inclusion contents of

the welds were estimated at 17 000 and 11 000 inclusions per squared

millimeter for the SN and DN welds respectively. Inclusion counts were made

at 2000X using an image analysis system on a light microscope (10 fields per

weld were evaluated)

.

Mechanical Properties

The mechanical properties measured for the welds are given in Table 2.

The yield strengths of the SN (962 MPa) and DN (936 MPa) welds differ only

slightly. Although both welds had strengths of less than 1000 MPa, the

strengths are high enough to give a good indication of the toughness that

can be anticipated at 1000 MPa.

The fracture toughness values of the DN (132 MPa-m*1

) and SN (141

MPa-m^5 ) welds are typical of type 300 series stainless steel welds.

However, these 316LN welds were expected to increase in toughness with

decreasing inclusion content and this did not occur. The results show that

the SN weld, which has a higher inclusion than the DN weld, also has the

higher toughness of the two welds. To explore reasons for this anomalous

behavior, we examined the fracture surfaces of the weld specimens.

Table 2. Mechanical Properties of the Type 316L Welds.

Specimen

Strength

(MPa)

Yield Ultimate

Elongation

%E1

Fracture Toughness

(MPa-m^5 ) Kic(J)

DN -

1

936 1458 52 135

DN - 2 129

SN -

1

962 1352 24 140

SN - 2 142

Fractographv

The fracture surfaces of the tensile and compact tension specimens have

rib- like features formed by ductile failure at intercellular regions of the

microstructure. Figure 2 shows a portion of the fracture surface on the SN

compact tension specimen that is dominated by these rib- like features. This

fracture surface morphology, sometimes called corduroy fracture, has been

reported previously for type 316 [13], and for type 310 [14] stainless steel

welds. To gain a better understanding of how these features relate to the

weld microstructure, the fracture surface of this specimen was partially

242
Figure 2: The fracture surface of the SN compact tension specimen has riblike features formed during ductile failure of intercellular regions in the

weld microstructure.

Figure 3: A plane polished parallel to the fracture surface of the SN

compact tensile specimen shows void formation associated with the ferrite

phase (light contrast, linear features).

243
ground and polished for light microscopy. In Fig. 3, areas of weld

microstructure and regions of the fracture surface below the polished plane

are shown. We see that the ribbed fracture features (dark band of ductile

dimple rupture) run along the intercellular regions of the microstructure.

Ferrite (light phase) is present in these regions, and failure appears to

originate near the ferrite-austenite interface. Voids preferentially

nucleate at inclusions or precipitates near the ferrite-austenite interface,

and coalescence of these voids results in the ribbed features present on the

fracture surface.

A similar failure scenario was developed from observations made on

cross sections of the SN compact tension specimen as shown in Figure 4.

This specimen was sectioned perpendicular to the crack growth plane,

parallel to the crack growth direction. The regions near the crack tip,

where cracks are not fully linked, show the crack path generally follows

intercellular regions in the microstructure, identified by the presence of

the ferritic phase. Again voids are observed near the ferrite-austenite

interface and they appear to be nucleated at inclusions. The ferrite phase

dig not crack or behave in a brittle manner in these specimens, nor in cross

sections of the tensile specimens which had extensive plastic flow prior to

failure

.

Figure 4: On cross sections polished perpendicular to the fracture plane in

the compact tension specimens, the crack path was found to be associated

with the ferrite phase. The ferrite is visible here (not etched) as a

skeletal network that is only slightly darker than the austenite matrix.

244
DISCUSSION

We suspect that the toughness of the intercellular regions is what has

been measured in these tests and that the importance of average inclusion

content and strength play a diminished role in determining fracture

toughness when failure occurs at intercellular regions. In the 18Cr-20Ni

(FN-O) and 316L (UNi, FN-7) welds that were tested previously (Fig. 1), for

example, inclusion content was shown to be a principal factor in limiting

the toughness of stainless steel welds whether the primary phase in

solidification is ferrite or austenite. In these cases, where inclusions

affect toughness, fracture surfaces have a more transgranular appearance

than those of the 316LN (12Ni, FN-4) welds tested in this study. The

intercellular fracture of the 316LN and the lack of a significant effect of

inclusion content on toughness may be due to the weld solidification mode.

As shown in Fig. 5, the liquidus surface of the Fe-Cr-Ni equilibrium

diagram, the type 316 stainless steel weld compositions cross over the

eutectic trough. Therefore, large regions of these weld structures can have

eutectic phase morphologies (austenite + ferrite) and compositions. The

316LN welds made in this study contain appreciable amounts of eutectic

structure. The 316L welds previously tested showed little evidence of

eutectic ferrite and welds with similar Cr content to these welds, but

higher Ni solidify in a fully austenitic mode. Comparing the 316 data, we

find significant gain in the strength- toughness relationship due to a change

in the solidification mode from FA (316L welds, FN-7) to AF (316LN, FN-4).

There are significant changes in the structures of these two types of welds

(austenite morphology, ferrite morphology, ferrite composition,

segregation) , but the effect of these differences on the strength and

Figure 5: The liquidus surface of the Fe-Cr-Ni system. The specification

limits for type 316 stainless steel welding compositions are indicated by

the parallelogram at the liquidus trough. Also, the positions of 18Cr-20Ni

alloys and type 904 stainless steel alloys are shown.

Cr

Nickel, wt.%

245
toughness of the welds is not resolvable in our data. Quite possibly,

beneficial effects on toughness, such as slightly higher Ni and lower

ferrite content for the 316LN weld, are offset by the compositional

inhomogeneities inherent in its more eutectic -like solidification. Although

the toughness of the 316LN welds might be improved by optimizing Cr-Ni

contents, increases in toughness are expected to be small.

To improve toughness significantly, it appears that major increases in

Ni content are needed, increases beyond the specification limit for type 316

stainless steel weld compositions. Such increases in Ni move the alloy

composition away from the liquidus trough, so the benefits of Ni on

toughness can realized [2]. Therefore, to weld type 316LN base materials,

it seems that welds of matching composition cannot be used if the strength

and toughness of the base material must also be matched. The 18Cr-20Ni and

the type 904L weld compositions shown on the Fe-Cr-Ni ternary diagram, have

proved to have mechanical properties more similar to the type 316LN base

metal

.

CONCLUSIONS

(1) A lower toughness intercellular region in the welds inhibited

improvements in toughness through reduction of the inclusion content.

(2) The type 316LN weld compositions tested in this study do not have

high enough toughness at 4 K to match the properties of type 316LN base

materials

.

REFERENCES

1. J.L. Young, C.J. Heyna, P. Komarek, H. Krauth, G. Vecsey, and C.

Marinucci, The Forced Flow Cooled Coils For the International Energy

Agency Large Coil Task, Advances in Cryogenic Engineering, Vol. 27,

Plenum Press, 1981, pp 11-19.

2. N.J. Simon and R.P. Reed, Strength and Toughness of AISI 304 and 316 at

4 K, Journal of Nuclear Materials, 141-143 (1986), pp 44-48.

3. G.F. Vander Voort and H.M. James, Wrought Stainless Steels, Metals

Handbook, 9th Ed., Vol. 9, ASM, 1985, pp 279-304.

4. Metals Handbook, Ed. 9, Vol 9, American Society for Metals, pp 279-296.

5. N. Suutala, T. Takalo, and T. Moisio, Ferrite-Austenitic Solidification

Mode in Austenitic Stainless Steel Welds, Met Trans. A., Vol 11A, 1980,

pp 717-725.

6. J . A. Brooks, J.L. Williams, and A.W. Thompson, STEM, Analysis of

Primary Austenitic Solidified Stainless Steel Welds, Met. Trans. A.,

Vol. 14A, 1983, pp 23-31.

246
7 . Y. Arata, F. Matsuda, and S. Katayama, Solidification Crack

Susceptibility in Weld Metals of Fully Austenitic Stainless Steels

(Report II) - Effect of Ferrite, P, S, C, Si and Mn on Ductility

Properties of Solidification Brittleness, Transactions of JWRI , Vol. 6,

No. 1, Welding Research Institute, Osaka, Japan, 1977, pp 105-116.

8. R.L. Tobler, H.I. McHenry, and T.A. Siewert, Strength -Toughness

Relationship for Austenitic Stainless Steel Welds at 4 K, Cryogenics,

Vol. 26, 1986, pp 392-396.

9. C.N. McCowan and T.A. Siewert, Fracture Toughness of 316L Stainless

Steel Welds With Varying Inclusion Contents at 4 K, Advances in

Cryogenic Engineering (Materials), Vol. 136, Plenum Press, 1989, pp

1331-1338.

10. C.N. McCowan et. al, Inclusions and Fracture Toughness in Stainless

Steel Welds at 4 K, Advances in Cryogenic Engineering - Materials, Vol.

32, 1986, Plenum Press, pp 335-342.

11. C.N. McCowan and T.A. Siewert, Stainless Steel Weld Metal: Prediction

of Ferrite Content, Welding Research Council Bulletin 342, 1989, pp 1-

36.

12. R.L. Tobler, D.T. Read and R.P. Reed, Strength/Toughness Relationship

for Interstitially Strengthened AISI 304 Stainless Steels at 4 K,

Fracture Mechanics: 13th Conference, ASTM STP 743, ASTM, 1981, pp 250

-268.

13. J.C. Lippold, M.C. Judas, and E.N.C. Dalder, The Relationship between

Microstructure and Fracture Behavior of Fully Austenitic Type 316L

Weldments at 4.2 K, Metallurgical Transaction A, Vol. 16A, 1985, pp

1835-1848.

14. C.N. McCowan and T.A. Siewert, The Fracture Toughness of 25Cr-22Ni-4Mn2Mo Stainless Steel Welds at 4 K, NISTIR 89-3931, Material Studies for

Magnetic Fusion Energy Applications at Low Temperature XII, 1989, pp

187-198.

15. Metals Handbook, Ed. 8, Vol. 8, American Society for Metals, pp 424.

247
'

10^
TEST STANDARDS

249

UNITED STATES -JAPAN DEVELOPMENT OF TEST STANDARDS PROGRAM

I

LEADER: R.L. Tobler

CONTRIBUTORS: P.T. Purtscher, R.P. Reed, D.P. Vigliotti, and R.P. Walsh

OBJECTIVES

o Draft mechanical property test standards for structural

alloys and composites at 4 K.

o Participate in other standards organizations including the American

Society for Testing and Materials (ASTM) and its task group activities

for the purpose of establishing the documents as voluntary concensus

standards

.

o Conduct experimental research to evaluate technical questions relating

to standard test procedures at 4 K.

o Administer interlaboratory test programs to support standards

developments

.

PROGRAM HIGHLIGHTS:

Structural Alloys

o Nine drafts of a proposed standard 4-K tensile test procedure were

written and submitted for review at domestic and international workshops. The ninth draft has been formally submitted to ASTM, and is

two -thirds of the way through the balloting process required to

establish a new standard.

o At ASTM's request the sixth draft of a proposed standard for 4-K

fracture toughness testing was rewritten in the form of an annex for

ASTM Method E 813-89; the annex was submitted to ASTM to begin due

process for standardization.

o Interlaboratory tension and fracture toughness tests were conducted on

a Japanese Fe-24Cr-15Ni alloy. A summary of results was published in

the proceedings of the International Cryogenic Materials Conference

held July 24-28, 1989 in Los Angeles.

o Research on the problems of Charpy impact testing at 4 K was conducted.

NIST, the Massachusetts Institute of Technology, and the Japan Atomic

Energy Research Institute cooperated to produce three research papers

on this subject. On the basis of this work, balloting has begun at

ASTM on a revision of the notched bar impact test standard, Method E

23-88.

o Standards development work for structural metals is now sufficiently

advanced at ASTM and prestandards research for composites has begun.

251
Composites

o Various national and international meetings were attended to establish

liaison with task groups involved in standards development. These

included:

1. ASTM's biannual D-30 technical committee meetings,

2. SACMA's Interface '89 conference on implementing composite

standardization,

3. the National Planning Committee's meeting to achieve industrywide standardization for composite materials, and

4. VAMAS' Technical Working Group meeting on Cryogenic Structural

Materials

.

o Composite materials scientists were interviewed to obtain information

necessary to identify the status of activities in standards development

for shear and compression tests of composites.

o An international symposium on composites testing was organized by NIST

as part of the International Cryogenic Materials Conference in Los

Angeles, CA, July 24-28, 1989. Special sessions were devoted to shear,

compression, and irradiation testing. The 16 papers presented were

published along with a summary paper to establish a point of departure

for initiating research in support of cryogenic standards.

o Improved experimental apparatus was constructed at NIST to evaluate

optimum test procedures at temperatures as low as 4 K. Interchangeable

fixtures allow various specimen types to be tested in the same

cryostat. Initial evaluation of shear testing has begun with

comparison of torsion, short beam, Iosipescu, and double -notched

compact specimen shear tests of glass-epoxy composites.

252
VAMAS Technical Working Party

Business Meeting, March 1990

CRYOGENIC STRUCTURAL MATERIALS

VAMAS (the Versailles Project on Advanced Materials and Standards) is

an international organization that conducts prestandards research to promote

and accelerate the development of standards for advanced materials. On

March 5 and 6, representatives from Austria, England, Germany, Italy, Japan,

Switzerland, and the U.S. met to discuss the VAMAS projects on cryogenic

structural materials and test standards. Kemforschungzentrum, Karlsruhe,

West Germany was the location. The first day (March 5) was devoted to

planning current programs up to the end of 1991. The second day explored

possible work beyond 1991.

First Dav

Host Dr. A. Nyilas delivered the opening remarks and welcome. Dr. T.

Ogata then presented the goals of the technical working party and outlined

the current status of round robin tests for alloys. Specific technical

presentations followed by Drs. Reed (standardization issues), Nyilas

(fatigue crack growth), and Ferrero (strain gage research). Presentations

were mixed with questions and discussion. The following is a summary of

information related to ongoing cooperative works.

The subject receiving most attention was the forthcoming round robin on

stainless steel. Seven labs in Japan and four in Europe and the U.S. have

agreed to participate. The test material is modified A-286 (Fe base, 26 Ni,

14Cr, 2Mo, 2Ti, and O.lAl), a potential candidate for superconducting

conduit sheaths with Nb 3 Sn conductors. Five tension and five toughness

specimens will be distributed to each laboratory in May. Several points of

protocol were resolved at this meeting:

(1) tension and toughness tests will be carried out using

the draft documents that NIST has written for ASTM,

(2) selected interpretive tests will be carried out at certain

institutes to supplement the basic program, and

(3) the test material will be properly screened beforehand,

by tests at NRIM.

The supplementary tests will be performed at various institutes and are

intended to answer remaining questions such as load cell calibration and

specimen size effects.

The 4-K strain gage calibration program was discussed by C. Ferrero.

In the past few weeks, materials and strain gages from major European, U.S.,

and Japanese manufacturers were distributed to all participants. This round

robin program is expected to take about six months to complete. Ferrero

informed us about the thermal expansion and Kondo resistivity effects

expected to be present in the calibration of Cu-Ni alloy strain gage outputs

on this program.

253
Second Dav

This was an informal meeting (with lower attendance) to explore work

beyond 1991. Considerable time was spent discussing a possible 5-year

effort devoted to composites. With draft standards for metals and alloys

being processed in several countries, it appears that composites pose the

next crucial problems in structural materials. A VAMAS program is desirable

because no single laboratory has the resources to tackle this subject alone.

The testing of composites is not a new idea; it was proposed at

previous VAMAS working meetings, and a special session was held in July 1989

at I CMC to define the state of the art in this area. Many different tests

are used for structural insulating composites in different countries, and

results for bulk materials are not in agreement, even at room temperature.

For example, shear testing at 76 and 4 K yields ultimate strengths that

range from 8 to 160 MPa for similar materials. This wide disparity is

frustrating and clearly unacceptable for designers. It is unclear how the

results for different test methods might be related to each other, and it is

impossible to specify the appropriate values for design calculations.

Specifically, a program to promote standard cryogenic shear test

methods for filamentary polymer-based composites was strongly advocated by

R.P. Reed. After initial discussions, the attendees at this meeting

focussed on the possibility of initiating a program with Dave Evans

(Rutherford National Laboratory) as the likely leader for such work. The

material of primary interest would be a glass—reinforced epoxy, perhaps

G10CR. First, a large batch of commercial material would be obtained and

characterized thoroughly with respect to composition and microstructure.

Then a double round robin would be executed in three steps:

I. Conduct the first round robin, showing the spread of results

for various shear test methodologies,

II. Draft the proposed shear test standard, and

III. Conduct a second round of tests to refine the draft.

While this plan progresses, volunteer labs would conduct selected

supportive research. The results would be assembled and analyzed to find

correlations between different techniques. A large reserve of material

would be available for supplemental or future use so that a large database

would continuously build for at least one material. The results could lead

to the development of miniature specimens for cryogenic irradiation studies.

In the past, this task had been tabled due to lack of a lead laboratory

with composites expertise, insufficient number of qualified labs (at least

six are required for round robins), and the priority of establishing metals

standards. At this time, Dave Evans can provide the necessary leadership,

and Rutherford has already provided his initial support for material

procurement and characterization expenses. The number of participating

laboratories would be increased to a sufficient level by inviting new

participation from organizations known to be interested in composites at

cryogenic temperatures. (A list of prospective organizations was readily

assembled at this meeting from the attendee's recommendations.)

254
SIDE GROOVES AND KIc CALCULATIONS*

R.L. Tobler

Materials Reliability Division

National Institute of Standards and Technology

Boulder, Colorado

This note discusses the method of stress intensity factor (K)

calculations for side grooved specimens of a high strength

austenitic steel. Two specimens with grooves and two without

grooves were tested. The Kj c measurements are equivalent if

calculations for grooved specimens use the net thickness instead of

the gross or effective thicknesses.

INTRODUCTION

Side grooves are sometimes advocated for ductile fracture tests

according to ASTM Methods E 813-87 [1], and yet no provision is made for

their use in linear elastic tests according to ASTM Methods E 399-83 [2].

Given a material of unknown toughness, it may be desirable to side groove a specimen, test it, and then (depending on the loading response) evaluate the

toughness using either the Kj c or the Jj c test methodologies. Permitting

side grooves in tests of linear elastic materials would facilitate the

development of a common specimen geometry and test procedure for Kj c and Jj c tests

.

For standard compact specimens without side grooves,

K - (P/(BW*5 )}-f(a/W), (1)

where

P

B

W

a

f(a/V)

applied load,

gross thickness,

specimen width,

average crack length, and

a function of relative crack length [2].

255
With side grooves, the thickness in the fracture plane is reduced and

Eq. (1) must be modified. Two substitutions for B in Eq. (1) are plausible.

Methods E 813-87 Section A2.5.3 [1] substitutes an effective thickness,

(B*Bn )

11

, where Bn is the net thickness. Freed [3], however, simply

substituted Bn

in his study of the effects of side grooves on single -edge - notched specimens of various alloys at room temperature. The more accurate

procedure can be determined experimentally. Here, we compare measurements

for plane (ungrooved) and side grooved compact specimens of an austenitic

steel tested in liquid helium at 4 K. The plane specimens were tested

previously by Read and Reed [4]. Two additional specimens remaining from

their study were fatigue cracked, then side grooved and tested.

PROCEDURE

The test material was a 5 cm thick steel plate in the as -received

condition: annealed at 1120°C for 2 h and water -quenched. The mill chemical

composition in mass per cent is Fe-21.15Cr-12.37Ni-4.96Mn-0.31N-0.041C2 . 17Mo-0 . 49Si- 0 . 015S -0 . 026P-0 . 18Nb-0 . 15V . The alloy in this condition has an

austenitic microstructure, an average grain diameter of 34 /im, and some grain

boundary precipitation [4]

.

Tension test results for this alloy are listed

in Table 1.

Fracture toughness tests followed Methods E 399-83 and used compact

specimens 38.1 mm thick and 76.2 mm wide. The crack orientation was T-L.

The specimens were not tested at room temperature, but at liquid helium

temperature (4 K) where linear elastic behavior prevails. Essentially, the

Kj c test at 4 K does not differ much from that at 295 K, except for the

provision of cryogenic apparatus.

The specimen geometry used in these tests [4] deviates from Method E

399-83 only in that the crack starter notch was modified to enable the clip

gage to be positioned at the loadline. However, a finite element analysis

[5] indicated that such modifications have negligible effect on K

calculations for relative crack lengths (a/W) greater than 0.4.

There was a slight difference in specimen preparation procedures in that

the specimens of the original study were precracked at 76 K using constant

minimum and maximum fatigue loads; the maximum fatigue stress intensity

factor steadily increased with crack length, reaching 35 MPa-m^ at the end of

Table 1. Mechanical properties of an Fe-21Cr- 12Ni-5Mn-0 . 31N steel (two

tests per temperature, averaged).

Temperature Yield Strength Ultimate Strength Elongation Reduction

of area

K MPa MPa % %

295 528 842 35 61

4 1446 1893 9 33

256
precracking. In the present study, the test specimens were fatigue cracked

in liquid nitrogen at 76 K, side grooved, and then fractured at 4 K at a

crosshead rate of 0.03 mm/s . The fatigue cycling was computer-controlled

with a sinusoidal wave in load control at 30 Hz. The minimum- to -maximum load

ratio during fatigue was 0.1. The stress intensity factor range AK in

fatigue was held constant at 35 MPa-m^ by load reductions as a function of

crack growth. The side

was 30.50 mm, and (B*Bn ) were 0.25 mm and 60°, respectively

grooves were 3.8 mm (10% of B) deep on each side, B

was 33.3 mm. The groove radii and included angles

n

RESULTS

A 4-K test record for a Fe-21Cr-12Ni-5Mn-0. 31N specimen with side

grooves is shown in Fig. 1. Type 1 load-displacement behavior is observed,

with slight nonlinearity before fast fracture. The critical loads Pq used

for calculating Kj c from Eq. (1) were determined by the secant offset

procedure in the usual manner [2].

Table 2 lists the Kj c values and validity criteria for the side grooved

specimens. The net specimen thickness is more than twice that is required

for valid tests. If Bn is used instead of B in Eq. (1), the toughness

calculations for side grooved specimens agree with the results for plane

specimens tested by Read and Reed [4], If (B'Bj^)^ is used, the calculated

fracture toughness values for side grooved specimens are 10% lower than those

for the plane specimens.

A representative 4-K fracture surface is shown in Fig. 2. Whether side

grooved or not, the Fe-21Cr- 12Ni-5Mn-0 . 31N steel specimens after testing

show essentially zero slant fracture along the edges. The fatigue crack

curvature is pronounced, but not drastic enough to disqualify the test

results. The edge cracks ae are 12-4% shorter than the 3-point average crack

lengths aa , but up to 15% is tolerable [2]. In this regard, note that Read

and Reed's datum [4] for one ungrooved specimen was reported as invalid (Kq)

in 1979, because the crack front curvature was excessive according to the

standard of those days (Method E 399-74). Today, that Kq value qualifies as

a valid Kj c measurement because the curvature tolerances were relaxed in the

updated standard [2]. The degree of crack curvature was further alleviated

by side grooving after fatigue cracking because that removes material from

the specimen edges and tends to reduce the difference between ae and aa.

DISCUSSION

At 4 K, the Fe-21Cr-12Ni-5Mn-0. 31N steel is linear elastic because of

its high nitrogen content and less than optimum metallurgical condition. The

precipitates in the microstructure are conducive to linear elastic fracture

and are not usually desired for cryogenic service. The specimen thickness,

whether side grooved or not, is more than sufficient to establish plane

strain in this study.

257
Figure 1. Load—displacement record for a side grooved specimen

tested at 4 K.

Table 2. Toughness results at 4 K for side grooved compact specimens, with

toughness calculated using different corrections for thickness.

The values for ungrooved specimens were 104 and 108 MPa'm*5 [4]

.

Specimen

number

Pm/Pq aazae aa

mm

Required

Thickness 3

mm

Kic

(using B-Bn)*1 MPa • in

4

KIc

(using Bn) MPa * in

5

2 1.05 12 13.7 95 107

12 1.03 14 13.4 94 106

Calculated from B - 2 . 5 (Kj c/YS)

2

, as specified by E 399-83, Section 9.1.3.

258
Figure 2. Appearance of a 38 nun thick compact specimen with side

grooves after fracture at 4 K.

Side grooves are sometimes used in ductile fracture tests for increased

constraint, but they cost more to machine and are not needed if a high degree

of plane strain constraint is already achieved in the gross thickness. For

ductile alloys, several benefits may derive from side grooves: they suppress

the formation of shear lips at the specimen surfaces in plane stress, improve

the uniformity of the stress state through the thickness, improve compliance

predictions of the crack length, and tend to alleviate tunnelling of the

crack front. Side grooves may also help to define the point at which slowstable crack extension begins in J-tests [6]. Side grooves are of little

practical value for linear elastic specimens but there is no reason to

exclude the use of side grooves in Kj c tests if that serves a purpose such as

the development of a common test method.

259
CONCLUSION

If side grooves are used in linear elastic tests the most accurate

method of calculation is desired. For compact specimens of a high strength

austenitic steel in this study we obtained best results by substituting Bn

for B in the conventional K-calibration , Eq . (1). Use of the effective

thickness (B-Bn

)^ appears to be less accurate, giving toughness values that

are 10% lower than Kj c. The experimental results favor the procedure of

Freed: Bn replaces B in Eq . (1). The implication is that the elastic strain

energy works uniformly across the net thickness, and that any perturbation of

the elastic strain energy at the specimen edges caused by the presence of

side grooves is negligible for thick specimens. Our experiments, however,

are limited to only a few tests of a material that exhibits a fairly high

degree of crack front curvature. More tests of other alloys having

straighter crack fronts are necessary to confirm this study.

REFERENCES

1. Standard Test Method for Jjc» A Measure of Fracture Toughness, ASTM

Designation E 813-87, 1988 Annual Book of ASTM Standards, Section

3, Metals Test Methods and Analytical Procedures (Amer. Soc. Test.

Maters., Philadelphia, 1988) 686-700.

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

Materials, ASTM Designation E 399-83, 1988 Annual Book of ASTM

Standards, Section 3, Metals Test Methods and Analytical Procedures

(Amer. Soc. Test. Maters., Philadelphia, 1988) 480-504.

3. C.N. Freed, Application of side grooves in the determination of Kj c for

high-strength metals, in: Fracture Mechanics of High Strength

Materials: Theory and Practice, ISI Pub 120 (Iron and Steel Institute,

London, 1970) 29-34.

4. D.T. Read and R.P. Reed, Toughness, fatigue crack growth, and tensile

properties of three nitrogen- strengthened stainless steels at cryogenic

temperatures, in: The Metal Science of Stainless Steels, E.W. Collings

and H.W. King, eds . (TMS-AIME, New York, 1979) 92-121.

5. R.L. Tobler and W.C. Carpenter, A numerical and experimental

verification of compliance functions for compact specimens, Eng. Fract.

Mechs., 21: 547-556 (1985).

6. P.K. Poulose , Determination of fracture toughness from thin side -grooved

specimens, Eng. Fract. Mechs., 26: 203-211 (1987).

260
I NTERLABORATORY TENSION AND FRACTURE TOUGHNESS TEST RESULTS FOR

CSUS-JN1 (Fe-25Cr-15Ni-0.35N) AUSTENITIC STAINLESS STEEL AT 4 K*

H. Nakajima, K. Yoshida, and S. Shimamoco

Japan Atomic Energy Research Institute

Naka-machi, Ibaraki, Japan

R.L. Tobler, R.P. Reed, R.P. Walsh, and P.T. Purtscher^

National Institute of Standards and Technology

Boulder, Colorado, U.S.A.

ABSTRACT

Interlaboratory tests are part of the U.S. -Japan cooperative program in

fusion energy to establish cryogenic test standards for structural alloys.

The second round of 4-K tension and fracture toughness tests for CSUS-JN1

(Fe-25Cr- 15Ni-0 . 35N) austenitic stainless steel are described in this

paper. The scatter of interlaboratory measurements is acceptable if some fracture toughness data are excluded as outliers.

INTRODUCTION

A data base and a mechanical design standard are required to build

superconducting coils for fusion reactors. The Japan Atomic Energy

Research Institute (JAERI) in collaboration with Japanese steel industries 1 successfully developed Japanese Cryogenic Steels (JCS) that have 4-K yield

strengths over 1,200 MPa, and fracture toughnesses over 200 MPaVra. Still,

the design data base is insufficient and there is no design standard yet

because the collaboration of many organizations is necessary to establish

them.

The U.S. —Japan Workshop on Low Temperature Structural Materials and

Standards was held in December, 1984. The ultimate goals of this work are

to establish the data base and design standard. The first collaborative

program was initiated to evaluate existing material test procedures and

standardize them. For this purpose, the first interlaboratory tension and

fracture tests for CSUS-JK2 (Fe- 22Mn- 13Cr- 5Ni-0 . 2N) austenitic stainless

steel were conducted at 4 K in 1987. 2 Drafts of the cryogenic test

standards were written based on those results. This paper presents the

second interlaboratory test results for CSUS-JN1 (Fe-25Cr-13Ni-0. 35N)

austenitic stainless steel that was tested according to the drafted

standards

.

Contribution in part by NIST; not subject to copyright in the U.S.

^•Present affiliation: Advanced Steel Processing and Products Research

Center, Colorado School of Mines, Golden, Colorado.

Advances in Cryogenic Engineering (Materialst. Vol. 36

Edited by R. P Reed and F. R Fickeu Plenum Press, Neu- York. 1990

1069
Table 1. Participating Laboratories and Personnel

Institute Supervisor Researcher Staff

(a Hitachi Research T. Matsumoto Y. Wadayama H. Sato

Laboratory (HRL)

(a Japan Atomic Energy S . Shimamoto H. Nakaj ima M. Oshikiri

Research Institute (JAERI) K. Yoshida

(a Kawasaki Steel Corp. K. Nohara N. Matsuno T. Katoh

(KSC)

(3 Kobe Steel, Ltd. (KSL) T. Horiuchi M. Shimada S . Nakayasu

<a Lawrence Berkeley

Laboratory (LBL)

J.W. Morris, Jr. J.U. Chan

J. Glazer

P.A. Kramer

(3 Lawrence Livermore J.R. Miller L.T. Summers D.A. Freeman

National Laboratory (LLNL) R . A . Riddle

(3 Massachusetts Institute

of Technology (MIT)

R.G. Ballinger I . S . Hwang

M.M. Morra

(3 National Institute of

Standards and Technology

R.P. Reed R.L. Tobler

P.T. Purtscher

R.P. Walsh

(NIST)

@ NKK Corp. (NKK) Y. Kohsaka N. Yamagami

@ National Research

Institute of Metals

(NRIM)

K. Ishikawa T. Ogata

K. Nagai

0. Umezawa

T. Yuri

(3 Nippon Steel Corp. (NSC) H. Abo T. Nakazawa

T. Takeshita

K. Oki

(§ Tohoku University, H. Takahaski T. Shoj

i

K. Nakano

Research Institute for

Strength and Fracture

of Materials (RISFM)

(3 Toshiba Research and H. Ogiwara E. Fukushima S. Kobatake

Development Center (TRDC) M. Tanaka

PURPOSE AND PARTICIPANTS OF THE SECOND INTERLABORATORY TESTS

The tests described here were conducted according to the proposed

draft cryogenic standards that were based on room temperature standards

(ASTM E 8, E 8M, JIS Z 2241) and the first interlaboratory test results.

Several aspects of testing at 4 K, such as test rate, specimen size, and

serrations, are considered in those drafts. The purpose of the second

round tests is to confirm the procedures defined by the draft cryogenic

standards and to check some items which must be resolved before the draft

standards are ratified.

The participants are volunteers from government, academic, and

industrial organizations with recognized experience in materials testing.

Table 1 lists the laboratories and personnel. Thirteen laboratories

participated, including nine from Japan and four from the U.S.

MATERIAL AND SPECIMENS

The test material is CSUS-JN1, one of the new JCS austenitic stainless

steels. It was supplied from a commercial steel manufacturer as 100 mm thick plate from a 50 metric ton industrial heat. Table 2 lists the

chemical composition.

262
Table 2. Chemical Composition of CSUS-JN1

c Si Mn P S Ni Cr N

0.018 0.33 4.1 0.019 0.002 15.26 25.2 0.37

Samples were cut at one laboratory from 1/4 thickness locations of the

original plate stock. This was done to minimize variables such as machining and specimen location as factors influencing the results. We

selected 18-mm diameter bars for tension tests, and 25-mra thick compact

tension (CT) specimens for fracture toughness tests, and distributed them

to each participant. The tension specimens were machined to their final

dimensions by each laboratory to suit individual preferences.

Tensile Specimens

The tension specimens were machined in the T orientation (load axis is

perpendicular to the hot-rolled direction) and in a cylindrical bar

configuration. Table 3 lists the final specimen dimensions for each

laboratory

.

Fracture Specimens

Most specimens used in this study were proportional CT specimens

(defined by Method E 813-81) with a 25 mm thickness, a 26.5 mm notch

length, and a 50 mm width. The CT specimens used for multiple-specimen

tests were proportional but were 20 mm wide. The notch orientation was TL, and displacement was measured at the loadline. Fatigue precracking

varied for each laboratory. Maximum stress intensity factors ranged from

22 to 43 MPa/ra, and initial crack- length- to -width ratios were 0.55 to 0.70.

Most specimens were precracked at room temperature (RT) , but specimens F8

,

9, 14, 15, and 16 were precracked at liquid nitrogen temperature (LNT) , and

F14, 15, and 16 were side-grooved after precracking to a net thickness

reduction of 20%. Except for side grooving, these variations in specimen

preparation did not lead to any resolvable effects on the fracture

toughness measurements.

Table 3. Nominal Specimen Dimensions (mm)

Specimen D G A R L

D: Diameter

T1.T2 6.0 30.0 36.0 15.0 85.3

G: Gage length

A: Length of reduced

T3.T4 7.0 35.0 .42.0 20.0 105.0 section

T5,T6 7.0 35.0 42.0 20.0 105.0 R: Radius of fillet

T7,T8 5 .

0

a 25.0 30.0 15.0 100.0 L: Total length of

T9.T10 7.0 42.0 20.0 105.0 specimen

Til ,T12 6.3 25.4 34.3 4.75 77.7 (See ASTM E 8 or 8M, Fig. 8)

T13 6.3 25.4 33.3 4.75 77.7

T14.T15 6.3 25.4 31.8 4.76 63.5

T16.T17.T18 6.3 25.4 44.0 9.52 71 . 8 a: Classified as a subT19.T20 7.0 35.0 50.0 20.0 105.0 size according to

T21 , T22 6.25 32.0 35.0 10.0 83.0 the 4-K draft standard

T23.T24 7.0 43.0 20.0 105.0

263
TEST PROCEDURE

Tension Test

The tensile properties of interest are the yield strength (YS), the

ultimate tensile strength (UTS), the elongation (EL), and the reduction of area (RA) . YS is defined by the 0.2% plastic strain offset method applied

to load-strain gage, load-extensometer , or load—time chart recordings. UTS

is the maximum load divided by the original area. EL is computed by three

methods: dividing the change in gage length by the original gage length

(ELC), dividing the change of total length by the original length of

reduced section (ELL), and dividing the change in length of the reduced

section by the original length of reduced section (ELA) . RA is the change

in cross-section area divided by the original area. Most tests were conducted using screw-driven machines with

3 crosshead displacement control

and nominal strain rates less than 1 x 10" s* , where the nominal strain

rate is the crosshead speed divided by the initial length of the reduced

section. Young's modulus (E) measurements were optional.

Fracture Properties

The fracture properties of interest are Jj c and Kj c (J), as estimated

from Kj c (J) - Jj c

* E. ^Ic *- s defined by the intersection of the blunting

and linear regression lines as per ASTM Method E 813-81. Ten laboratories

used the single-specimen unloading compliance (ULC) technique, but the

multiple-specimen (MS) method and a special single-specimen Key-Curve

technique 2 were also used. Seven labs used servohydraulic machines (four

used clip gage control, three used stroke control). Five labs used screw- driven machines (crosshead speed control). Four labs used specimen

rotation and/or crack growth corrections to J—Aa data.

RESULTS

Tensile Properties

Figure 1 shows the YS and UTS data plotted versus nominal strain rate.

There are no measurable effects for the strain rates covered in this study:

0.23 to 8.3 x 10* 4 for YS, and 0.79 to 12 x 10* 4 for UTS. The time chart

YS data are nearly equivalent or slightly higher than results from strain

or extensometer plots, with differences ranging from -3.4 to 13.1%. Three

laboratories discovered an unusual serration phenomenon: at low strain

rates, as shown in Fig. 2, only one serration occurred before failure.

70

-60

§50

”40

2 30

3 20

10

0 10 20 30 40 50 60

TIME (min)

Fig. 1 Yield and ultimate strength

measurements at 4 K

Fig. 2 Load- time chart for a test at low strain rate

1072

264
Table 4. Tensile Properties of CSUS-JN1

Spec

.

YS a S-S

(MPa)

YS b

L-T

(%)

UTS

(MPa)

ELG

(%)

EL

ELL

(%)

ELA

(%)

RA

(%)

E

(GPa)

Strain Rate

(10’« s* 1 )

T1 1778 20 21 21 39 0.23, 1 .

2

C T2 1328 -0.3 1789 23 23 21 42 195 0.23, 1 .

2

C

T3 1356 1.0 1740 31 45 189 2.4, 12.

0

C T4 1371 2.1 1768 35 47 208 2.4, 12.

0

C T5 1426 1.8 1851 24 21 39 210 1.2, 0.79°

T6 1375 2.8 1893 25 24 44 194 1.2, 0 . 79 c T7 1384 -3.4 1876 37 37 35 43 240 6.7, 10.

0

C T8 1893 33 37 32 45 6.7, 10.0°

T9 1254 2.1 1770 33 45 191 4.0, 10.

0

C T10 1757 34 46 o *—

o o

o Til 1352 -0.5 1760 32 38 37 46 196 5.9

T12 1407 l

H* OO 1794 28 32 28 46 194 5.9

T13 1333 1872 39 45 208 8.3

T14 1379 1848 35 48 209 8 .

3

T15 1768 34 37 50 5.0

T16 1125 13.1 1855 24 22 17 38 195 0.9

T17 1330 5.5 1873 21 20 16 42 202 0.9

T18 1361 4.2 1845 24 20 15 37 202 0.9

T19 1787 38 42 38 46 6.7

T20 1754 38 43 33 52 6.7

T21 1822 31 32 42 210 4.8, 9 . 5 C T22 1263 8.7 1792- 37 35 34 208 4.8, 9 .

5

C T23 - - 1817 - - 29 29 d 47 2.0

T24 1797 27 27 d 49 - - 2.0

Max

.

(%) 6.7 4.5 9 13 10 8 Min. (%) -15.8 -4.0 -10 -10 -13 -10

Avg. 1336 1812 30 30 28 44

a: S-S •- Measured from stress-strain curve.

b: L-T -- Measured from load- time curve and reported as % of YS a

. c: A rate change was used; cited first is the rate before the YS

measurement; second is the rate after YS measurement until failure,

d: Measured from length of fillet.

A summary of tension test results is given in Table 4. The variation

in EL values obtained by the three different measurement methods is within

9% (see specimen T18) , but the scatter of measurements among laboratories

is about 20%. The relations between ELL, RA, and nominal strain rate are given in Fig. 3. At low strain rates, ELL increases with increasing

^ nominal strain rate and becomes constant at rates higher than 5 x 10" s' The effect of strain rate on RA is weaker compared to its effect on ELL.

Fracture Toughness

The fracture toughness data are listed in Table 5. Except for the

data of specimens F8 and 9 which indicate low toughness, and F10 and 11

which indicate high toughness, the bulk of the results are consistent as explained in the Discussion.

1073

265
Table 5. Fracture Properties of CSUS-JN1

Spec

.

J I C 2 (kJ/m ) Kic(J)

(MPa/m)

Cl a (MPa)

C2 a (kJ/m

2

)

a ob

(%)

Aa

p

b

(%)

Method Precrack Temp.

FI 254 226 235 234 0.2 -17.8 ULC RT

F2 243 221 199 227 -1.7 -0.6 ULC RT

F3 240 220 187 226 -3.0 -11.3 ULC RT

F4 224 213 173 211 -2.7 -15.3 ULC RT

F5 228 214 188 214 0.01 -4.1 ULC RT

F6 238 219 140 227 ULC RT

F7 208 205 270 190 -3.4 -10.0 ULC RT

F8 134 164 206 123 -4 -22 ULC LNT

F9 49 100 206 46 -8 0.7 ULC LNT

F10 466 305 195 439 ULC RT

FI 1 411 286 174 390 ULC RT

FI 2 216 219 132 207 0.3 24.0 ULC RT

FI 3 239 230 193 225 -0.5 12.0 ULC RT

F14 c 149 173 230 138 -0.03 -2.4 ULC LNT

F15 c 234 217 197 220 1.0 -10.9 ULC LNT

F16 c 163 182 178 153 -2 -8.9 ULC LNT

FI 7 207 204 269 189 -0.5 -36.1 ULC RT

F18 237 219 281 216 -0.6 -8.6 ULC RT

F19 231 216 Key -Curve RT

F20 - 24 d 254 226 270 232 MS LNT

Max . ( %

)

102 43

Min . ( %

)

-79 -53

Avg. 231 213

a: Cl and C2 are constants of the R-curve [J - Cl*Aa + C2].

b Calculated from (prediction measurement) /measurement

.

c: Specimens were side-grooved after precracking,

d: Crack extension is the average of center three points.

DISCUSSION

Tension Test

The YS and UTS measurement consistency in this study is similar to the

first inter laboratory tes^ results . It appears that limiting the nominal

strain rate to 1 x 10' s* or less, as in the current draft standards, is

reasonable. This study indicates that a small bias may be introduced in YS

50

5

40 - 30

-j 20

w 10

0

5x10*

9

1<54

1(J

S 2i«Js NOMINAL STRAIN RATE (**')

5i1<J

9

1<S4

1<5 S 2i1<5

3 NOMINAL STRAIN RATE (s'

1 )

Fig. 3 Elongation and reduction of area measurements at 4 K

1074

266
measurements when the 0.2% offsec procedure is applied to time chart plots,

rather than strain gage or extensometer plots. This bias depends not only

on the test apparatus but also the technician's skill. For convenience,

however, the load-time chart method is quite useful in alloy screening or commercial production control where maximum precision is not required.

ELL and ELA measurements are specified in the draft tension test

standards based on the assumption that plastic elongation of the gripped

part is negligible compared with that of the specimen's reduced section.

ELL and ELA are easy to measure, and are potentially useful since they

eliminate the need for gage length marks which sometimes cause failure. As

there are only small differences between ELG , ELL, and ELA in this study we can employ ELL and ELA in addition to ELG at 4 K.

Unusual serration behavior explains the observed dependence of EL on strain rate. When tested at low strain rates, specimens of JN1 showed

lower ductility and fewer serrations. This slow strain rate effect

increases the scatter of interlaboratory elongation measurements, and if it

is taken into account the overall data scatter may be reduced to half that

shown in Table 4. This effect depends on the material tested, and further

experiments will be needed to fully evaluate it. Revised draft standards

should note this effect, and a lower limit of strain rate should perhaps be

selected to establish the data base for such materials.

Fracture Toughness

The scatter of reported Jj c values, as shown in Table 5, is very large

in this study. However, if -the F8 , 9, 10, and 11 data are eliminated, the

pooled average values of toughness for specimens without side-grooves are

J Tc - 232 kJ/m

2 +9.5%, -10.8%,

and

Kic(J) - 216 MPaVm +4.6%, -5.6%.

Then, as before, 2 there is little scatter. The pooled data include a datum

obtained from five specimens tested by the MS method (F20—24, Table 5). In

the MS method, thinner specimens were used, but the data are comparable as there is no size effect for 12.5 to 75 mm thicknesses in this alloy.

Evidently, side grooving reduces for this material. In a separate

study of CSUS-JN1, side-grooved specimens gave about 20% lower values than

those without grooves. If the data from specimens F14, 15, and 16 are increased by 20%, those results, too, agree with the pooled average value.

The data of F8, 9, 10, and 11 clearly fall outside the range of pooled

data. The resistance curves for these specimens contained no data in the

blunting line region, and only three to five points in the regression line

region, which contributes to scatter in determination of Jj c

. For example,

resistance curves at the onset of testing often show negative Aa. In the

cryogenic standard, we treat this by shifting the blunting line to the left on the Aa axis until there is a best fit to experimental data. If there

are too few data in the blunting region, the blunting line is not shifted

to the best position. We conclude chat more data in the blunting region as well as in the regression line region should be taken when the unloading

compliance technique is used.

As shown in Table 5, the Aap agreement is not good in this study. In

most cases the ULC-predicted values are smaller than those measured from

fracture surface. The measured and calculated values must agree to within

15% according to ASTM Method E 813-81. Our Jj c data, however, are not

sensitive to the Aap disagreement, as also indicated previously. 2 Although

1075

267
che Aap agreement affects the slope of the resistance curve, apparently it

has a negligible effect on Jj c

.

SUMMARY

Eleven laboratories measured tensile properties and twelve measured

Jj c for CSUS-JN1 (Fe-25Cr-15Ni-0. 35N) austenitic stainless steel according

to the proposed cryogenic test standards. Aside from the slow strain rate

effect on elongation which is peculiar to this material, there are few

systematic effects for tensile properties as long as tension tests are conducted according to the proposed standards. Therefore standardization

of the drafts is being pursued through JIS in Japan and ASTM in the U.S.

The unloading compliance method appears to be reliable for Jj c measurements at 4 K, but some technical details remain to be clarified.

Modifications of che current draft standard and further experiments may be

required to achieve a consensus standard. The results suggest that the

current Aap agreement requirement may be relaxed, and that more than the

minimum of 4 J—Aa data points should be obtained when using the resistance

curve method.

ACKNOWLEDGMENTS

The authors would like to thank Drs . S. Mori, M. Yoshikawa, and M.

Tanaka for their continuous encouragement of this work. The organizers of

this activity wish to express appreciation to all participants for their

contributions. Without such collaboration, the development of consensus

standards is impossible. Finally, special appreciation is due to Nippon

Steel Corporation for providing the test material.

REFERENCES

1. H. Nakajima, K. Yoshida, M. Oshikiri, Y. 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 for forced flow superconductors at 4 K, in: "Advances in

Cryogenic Engineering— Materials", Vol. 34, A.F. Clark and R.P.

Reed, eds., Plenum, New York (1988).

2. H. Nakajima, K. Yoshida, S. Shimamoto, R.L. Tobler, P.T. Purtscher, and

R.P. Reed, Round robin tensile and fracture test results for an Fe22Mn-13Cr-5Ni austenitic stainless steel at 4 K, in: "Advances in

Cryogenic Engineering-Materials", Vol. 34, A.F. Clark and R.P. Reed,

eds., Plenum, New York (1988).

3. K. Yoshida, H. Nakajima, M. Oshikiri, R.L. Tobler, S. Shimamoto, R. Miura, and J. Ishizaka, Mechanical tests of large specimens at 4 K: facilities and results, in: "Advances in Cryogenic Enineering—

Materials", Vol. 34, A.F. Clark and R.P. Reed, eds., Plenum, New

York (1988).

4. H. Nakajima, K. Yoshida, and S. Shimamoto, Effects of specimen size,

side grooves, and control on fracture toughness in the JCS,

presented at The Third U.S .-Japan Workshop on Low Temperature

Structural Materials and Standards . Tokai-mura, Japan, 1988.

1076

268
Authorized Reprmt 1 990 from Journal of Testing and Evaluation. July 1 990

Copynght American Society for TesOng and Maaanais, 1916 Race Street Philadelphia, PA 19103

TECHNICAL NOTE

Y. A. Rosenthal,

1 R. L. Tobler and P. T. Purtscher3

J,c Data Analysis Methods with a “Negative Crack Growth”

Correction Procedure

REFERENCE: Rosenthal. Y A . Tobler. R. L.. and Purtscher. P T "Jw Data Analysis Methods with a 'Negative Crack Growth' Correction Procedure.” Journal of Testing and Evaluation. JTEVA. Vol.

IN. No 4. July IWO. pp 301—304.

ABSTRACT: It often occurs in J-R testing that some of the initial crack extension (Ad) data points have anomalous negative values. ASTM Methods E 813-87 and E 813-81. however, do not specify any

procedures for treating such data We propose a procedure for data

analysis which utilizes the negative la values and is consistent with ASTM E M3-N7 The proposed correction procedure is illustrated using a data set for austenitic steel: also. A values obtained using

ASTM E 813-87 are compared to those obtained from its predecessor E 813-81 KEY WORDS: ASTM Method E 813-87. fracture toughness. 7-in- tegral. J-R testing, negative crack growth. /?-curve

ASTM Method E 813-87 covers the experimental determination of 7,.. which can be used as an engineering estimate of

fracture toughness near the initiation of slow stable crack growth

for metallic materials. The single-specimen elastic compliance

technique (as described in Section 8.4 of that standard) is commonly used. According to that procedure, crack extension (la)

measurements, inferred by partial unloading, are performed fre- quently throughout the loading of the specimen. It often occurs, due to friction (/]. to misalignment in the loading train [2|. to the electronic equipment balancing and zeroing, or perhaps to physical blunting behavior effects (J). that some of the initial inferred la data points have anomalous negative values. ASTM

E 813-87. however, does not discuss the problem of "negative

crack growth" or specify any procedure for treating such data.

In a recent interlaboratory program, most of the participants

observed some negative crack growth in tests of a tough austenitic

steel (4j. It is not possible to repeat each test which resulted in

Manuscript received 7/20/89; accepted for publication 1/11/90. Note—This paper is a contribution of the National Institute of Stan- dards and Technology and is not subject to copynght in the United States. 'On a sabbatical from the Nuclear Research Center—Negev. PO Box

9001. Beer Sheva 84190. Israel. 'Xational Institute of Standards and Technology, Matenals Reliability Division. Boulder. CO 80303-3328. 'Advanced Steel Processing and Products Research Center. Colorado

School of Mines. Golden. CO 80401.

negative la data (especially when environmental effects are also involved), because of the cost. In many laboratories dealing with

J-R testing, this problem is treated by eliminating the negative

la values, or by shifting the resistance curve to a new zero, such

as the most negative la point. Such treatment is arbitrary and

undesirable, since it depends on the operator's decision. Obviously it can also be a source of scatter when comparing inter- laboratory toughness data for the same material tested under the same conditions.

The purpose of this Technical Note is to propose a data analysis

procedure for 7k determination which explicitly recognizes and

utilizes the negative la values and is otherwise consistent with

ASTM E 813-87.

Data Analysis Procedure

The data analysis procedure is carried out using a personal

computer with an interactive computer program especially de- veloped for this purpose. A typical 7-versus-Aa data set may

include some initial negative la values that cannot be analyzed

according to the existing standard, because it demands the use

of a method of least squares to determine the linear regression

line:

ln(7) = In ( C, ) + C ln(Aa) ( 1

)

where C, and C; are constants. According to ASTM E 813-87. one is asked to fit only the data which fall within the exclusion

interval (Section 9.2.5). In extreme cases it is possible that some

negative la values will be included in the acceptable data points,

but Eq 1 cannot accept negative values for la. In any case, the exclusion interval is defined in relation to the blunting line, the true zero of which is questionable in the case of negative la

values. Therefore the operator usually resorts to shifting the la

values arbitrarily.

The details of a better procedure, proposed by the authors,

are shown in Fig. 1. The first step is to fit all data to a nonlinear

(power) function having an inverse form of Eq 1:

la = K,(J)K1 + K> (2) 0090-3973/90/0007-0301 $00 00

301

269
302 JOURNAL OF TESTING AND EVALUATION

Raw Aa, J dataset

i

Fit non-linear function:

*a = Ki (J)*2 + K3

Plot Aa(J) data and the curve fit.J

Shift Aa values:

Aa’ = Aa - K3

where K„ Ku and K, are constants. This function is fitted to all the data, including negative An. In the next step, K, is identified and used as the amount needed for shifting the data to a new

zero. Then, the entire /- versus-An plot (including the blunting

line, the exclusion lines, the division lines, and the offset line)

is shifted along the abscissa by an amount, Ku thus creating new

Aa' crack extension values. The data points for the final curve

fit are defined in the third step. According to Section 9.2.5 of

ASTM E 813-87, a regression line having the form of Eq 1 is fitted to the acceptable data, and JQ

is determined at the intersection point between the nonlinear regression line and the 0.2 mm offset line (Fig. 2). The J0

values, as calculated from the regression line and from the offset line at the point of intersec- tion, converge to within less than 0.1%, whereas 2% or less is required in ASTM E 813-87, Section 9.2.6.4. For the sake of comparison, we analyzed one set of data three ways. Figure 2 shows the original data analyzed with the pro- posed procedure. Figure 3 shows the same data set, analyzed

after shifting the Aa values to the most-negative value. The analysis was also repeated for this data set after eliminating all the

270
ROSENTHAL ET AL ON J* OATA ANALYSIS METHOOS 303

Aa, mm

FIG. 2—J-Aa curve for a lough austenitic stainless steel at 4 K (obtained

using a standard 25.4 mm thick compact specimen in a liquid helium

environment), which includes negative Aa values, as analyzed using the proposed procedure.

FIG. 3—Same data set as shown in Fig. 2 analyzed after shifting the zero to the most negative Aa value.

negative An values (Fig. 4). 7k values, calculated in these three

different ways, are shown in Table 1. The results vary from 385

to 480 N/mm, depending on the data reduction procedure, but

the proposed procedure generally gives higher values.

In addition to the problem of negative crack growth, we

wanted to assess the significance of the revised standard test

FIG 4—Same data set as shown in Fig. 2 analyzed after eliminating

the negative Aa values.

TABLE 1—Comparison of lk obtained after different treatments to negative Aa values.

Procedure /k , N/mm

(1) Using all the data and the proposed procedure (Fig. 2) 480

(2) Zero shifted to the most negative Aa value (Fig. 3) 385

(3) Eliminating the negative Aa values (Fig. 4) 475

methodology on /k measurements. We analyzed a data base of

results based on 42 J-R tests (mostly at 4 K) with austenitic

stainless steel specimens (with yield strengths ranging from 400

to 1700 MPa) using ASTM E 813-87 and its predecessor E 813-

81. K, determined from the proposed correction procedure (Fig.

1) was used as the amount of crack extension for shifting the Aa

values in both cases. The 7k values obtained using the new

method are plotted versus Jk values obtained from the previous

method’s analysis (Fig. 5). For most cases, /k obtained using the new method is higher

than Jk obtained using the old method, and the difference be- tween the two 7k values tends to increase as the toughness in- creases. This trend is not influenced strongly by our negative

crack growth correction, but results primarily from the different definitions of 7k used in the two methods.

Coaduioii

1 . The data analysis procedure proposed here enables /k de- termination even when minor testing problems produce negative

values of crack extension. This proposed data analysis procedure

eliminates arbitrary treatment with respect to negative crack ex- tension values and therefore enables a more meaningful comparison oiJ-R curve results, as obtained by different laboratories,

by excluding operator bias.

271
304 JOURNAL OF TESTING AND EVALUATION

FIG 5—

J

k obtained using ASTM E 813-87 ( with the proposed correction procedure) plotted versus J k from ASTM E 813-81 analysis. The

dashed line represents equal J,c values.

2. J* data obtained using ASTM E 813-87 may be considerably

higher than those obtained using its predecessor ASTM E 813- 81 . The difference may be substantial, especially for high tough ness materials.

Acknowledgments

The authors would like to acknowledge Dr. D. T. Read. Na- tional Institute of Standards and Technology, for discussions during the development of the data analysis procedure

References

( / 1 Voss. B.. "On the Problem of Negative Crack Growth and Load Relaxation in Single Specimen Partial Unloading Compliance Tests,

in Ductile Fracture Test Methods (Proceedings of a CSNI Workshop

Pans. 1-3 Dec 1982). Organization For Economic Cooperation anc Development. Nuclear Energy Development. Paris. 1983. pp 210- 219

[2] Bowman. R.. Antolovich. S D . and Brown. R C . "A Demon

Stratton of Problems Associated with Crack Closure Meusuremen

Techniques." Engineering Fracture Mechanics. Vol 31. No 4. 1088 pp. 703-712.

( J] Purtscher. P T.. “Absence of Stretch Zones in Austenitic Stainles Steels Fractured at Cryogenic Temperatures. Journal of Tesungam

Evaluation. Vol. 15. No 5. 1987. pp. 296-298.

(4| Nakapma. H.. Yoshida. K.. Shimamoto. S . Toblcr. R L Purtscher. P. T.. and Reed. R. P.. "Round Robin Tensile jnd Frac ture Test Results for an Fe-22Mn-13Cr-5Ni Austenitic Stainless Stee at 4 K." in Advances in Cryogenic Engineering. Vol 34. A F Clarl and R. P. Reed. Eds.. Plenum. New York. 1988. pp 241-249

272
CHARPY IMPACT TESTS OF SUPERCONDUCTOR CONDUIT

ALLOY WELDS NEAR ABSOLUTE ZERO

R.L. Tobler and D.P. Vigliotti

Materials Reliability Division

National Institute of Standards and Technology

Boulder, CO 80303

This note describes a technique used at the Japan Atomic

Energy Research Institute (JAERI) and now at the National

Institute of Standards and Technology (NIST) for conducting

Charpy impact tests in flowing helium at an initial test

temperature near 5 K.

INTRODUCTION

Impact tests at temperatures near absolute zero are controversial at

present because there is no standard method for conducting the tests.

Certain regulatory agencies require impact tests on structural materials for

service at liquid helium temperature, even though there is no scientific

consensus on the meaning of impact measurements near absolute zero.

We previously reviewed the problems of impact testing at temperatures as

low as 4 K [1]. ASTM Method E 23-88 [2], the existing standard for notched

bar impact tests, specifies that the specimen must be transferred from the

cooling bath through air to the machine and tested within 5 s. That

procedure does not work at temperatures below 77 K. To attain more extreme

cryogenic test temperatures it is necessary to insulate the specimen from

contact with air, but Method E 23-88 does not specify an appropriate

procedure

.

Incoloy 908 is a newly developed alloy for superconducting conduit

applications in fusion energy magnets at 4 K. Research at the Massachusetts

Institute of Technology (MIT) shows that this material resists degradation

during thermal aging treatments and has desirable thermal expansion behavior

for superconductor sheath applications. The 4-K tensile, fatigue, and

fracture toughness properties of this conduit alloy were measured and

reported previously [3]. The alloy was selected for use in the United States

Demonstration Poloidal Coil (USDPC) which will be installed and tested at

JAERI according to international agreements. Before the coil could be

installed, nonstandard impact tests for the conduit alloy were required by

Japanese law. We tested the conduit alloy welds and two austenitic stainless

steel base metals for comparison as described in this report.

273
PROCEDURE

MIT provided the test material, fabricated the Charpy test specimens,

and sent them to NIST to be tested. The test material was a gas- tungstenarc (GTA) welded plate of 2.35 mm thickness identical to that of the outer

conduit of the USDPC. Conduit walls are typically 2 mm thick, whereas the

conventional Charpy specimen has a 10 x 10 mm cross section. A 10 x 10 mm

cross section was achieved in this study by welding a stack of 2 mm thick

laminates. These specimens were tested in the manner described to satisfy

Japanese law.

laminates were cut from solution annealed plate and wet- ground to 2 mm

thickness; 5 such specimens were electron beam welded to create each

composite Charpy test specimen. The GTA welded central region of the

composite specimen, a length of 10 mm, was not bonded by electron beam

welding in order to avoid any heating effect. The composite was then

machined to the standard Charpy size: 10 x 10 x 55 mm. Each specimen was

notched in the weld at the center of specimen length. The specimens were

then heat treated in vacuum for 200 h at 650°C to precipitation harden the

alloy. The notch location was checked after heat treatment by etching each

specimen to be sure that it was centered in the weld metal.

The service environment for the conduit is liquid helium at 4 K. Method

E 23-88 is invalid for tests at 4 K because moisture condensation on the

specimen quickly warms it up to about 60 K as it is being transferred through

air to the test machine. Therefore we used a variant of the "flow" method to

obtain an initial test temperature near 4 K.

Originally, the flow method was introduced by Jin, Horwood, Morris, and

Zackay [4]

.

It was later modified, greatly simplified, and used in Japan

[5]. The system used recently at JAERI is adopted here. The specimen is

wrapped in a thin sheet of plastic, a tubular container, that is attached to

a liquid helium transfer line. The container is placed on the test machine

anvil, and liquid helium is pumped through the tube. After cooling at least

two minutes, and while helium is still flowing through the tube, the specimen

and container are fractured by release of the pendulum.

The plastic wrapping used in our tests was transparent overhead

projection sheets, 0.1 mm thick. The sheets are cut to 220 x 55 mm and

marked with inklines for specimen positioning as shown in Fig. 1A. Doublesided adhesive tape is applied to the sheet at symmetrical locations to

secure the specimen as it is positioned on the sheet. The plastic is then

folded and wrapped around the specimen, and the seam is sealed with

cellophane tape (Fig. IB). Finally, the specimen container is taped to the

helium transfer line and carefully positioned with the notch centered on the

test machine anvil.

The 360 J test machine used in this study is a conventional swinging

pendulum machine with a C- shaped pendulum. The angle of pendulum swing is

read from the machine ' s dial indicator and the reading in degrees is

converted to absorbed energy using conversion tables.

274
_L

10

1

50

Plastic

Container

Wrapped Specimen

Figure 1. Plastic film wrapper (A), cross sectional view of wrapped

specimen (B)

.

The dimensions are mm.

275
During a test, liquid helium is transferred from a 100 1 storage Dewar.

In these tests, helium gas was used to pressure the storage Dewar, and the

flow of pressurizing gas was consistently maintained at a value of 30

standard cubic feet per hour. For this rate of transfer, it takes 2 min to

cool the specimen. The minimum temperature that can be reached depends on

some details of the apparatus, especially the length of the metal tube that

extends beyond the vacuum- j acke ted portion of the transfer line. With our

existing transfer line about 25 mm of tube was exposed (that is, wrapped with

foam and tape, inserted in the the specimen wrapper, but not vacuum

insulated) , and the minimum temperature that could be reached at the center

of the Charpy specimens was 5.5 K.

The specimen cooling time and initial test temperature were verified by

calibrations using a conduit alloy weld specimen containing a thermocouple.

We selected a Chromel versus gold + 0.03% iron thermocouple for maximum

sensitivity at 4 K. Our thermocouple wire diameters were 0.305 and 0.127 mm,

respectively, and we bonded them using a thermocouple welder at low power.

We drilled a 1 . 6 mm hole along the calibration specimen's axis to the base of

the notch, and then inserted the thermocouple. After adding some fine copper

powder to ensure good thermal contact between the thermocouple bead and the

specimen, we filled and sealed the hole with thermally nonconducting varnish.

We calibrated the thermocouple following a procedure suggested by Sparks

and Powell [6]. They published detailed calibrations for various chromel

versus gold + 0.02% iron thermocouples and proposed that their reference data

could be used for similar thermocouples if adjusted satisfactorily with data

from spot calibrations. For spot calibrations, we immersed our thermocouple

in liquid nitrogen and liquid helium baths. The difference between our

output and that reported by Sparks and Powell at 4 K was only 210 /iV (a

variation of 4.5%). We shifted their reference curves by that much. Based

on calibrations with the instrumented specimen in a container on the machine

anvil, we achieved an initial test temperature of 5.5 K in flowing helium.

Temperature -versus -time curves confirmed that this temperature is completely

stable after 2 min.

The impact test apparatus is illustrated in Figs. 2—5. Shown is the

impact machine with the pendulum raised (Fig. 2) , a container on the anvil

with the specimen centered in the gap (Fig. 3), helium flowing through the

line (Fig. 4), and a close-up depicting the moment of impact (Fig. 5).

RESULTS

The Charpy test results are listed in Table 1. Four weld specimens were

tested, and the average absorbed energy is 32.9 J. The individual values

range from 30.5 to 34.3 J. One plastic container with no specimen inside was

tested to confirm that very little energy is absorbed in fracturing the

insulation at 5.5 K. The extra energy added by the container is negligible

at not more than a few tenths of one joule. Accordingly, no calibration

correction was made to the data listed in Table 1.

276
Figure 2. Overall view of the impact test machine.

277
Figure

3.

Encapsulated

specimen

on

the

anvil

at

the

beginning

of

a

test.

Inked

lines

on

the

plastic

container

are

used

to

center

the

specimen.
279

Figure

4.

Helium

transfer

through

the

specimen

container.
r* 

Figure

5.

The

pendulum

as

it

collides

with

the

specimen

container

at
Table 1. Charpy test results for specimens initially at 5.5 K

Material Specimen Charpy absorbed energy JAERI's result

No. J a f t- lbfb

J

908 Weld 1A 33.5 24.7246

( laminated) IB 34.3 25.2799

1C 30.5 22.5072

2G__ 33.2 24.7246

Avg - 32.9

Fe-19Cr-8Ni-4Mn 1 101.8 75.0591

2 97.1 71.6403

Avg - 99.4 93

AISI 316LN 1 240.0 176.9678

z 221.4 163.3087

Avg - 230.7 230.

6

C

a: J - 1.356 times reading in ft-lbf.

b: Value in ft-lbf is read from table after noting angle of pendulum swing,

c: Average value for three specimens tested using the flow method, and

same alloy heat.

281
In addition, we conducted Charpy tests in duplicate for two austenitic

stainless steel base metals: an Fe-19Cr-8Ni-4Mn-0. 21N alloy and AISI 316LN.

Standard specimens (10 x 10 x 55 mm) were machined from commercially annealed

25 mm thick plates. The average absorbed energies for these steels were 99.4

and 230.7 J, respectively, as shown in Table 1. The same two steels were

tested in triplicate at JAERI. Using their virtually identical flow method,

they reported average absorbed energies of 93.7 and 230.6 J, respectively.

The NIST and JAERI data are in fair agreement, which gives some assurance

about the reproducibility of measurements obtained using the flow method.

ACKNOWLEDGMENTS

The flow test method is being used in an ongoing cooperative study by

JAERI, MIT, and NIST. Special acknowledgments are due to H. Nakaj ima (JAERI)

for sharing the details of this test technique with us, and to M.M. Steeves

and I.S. Hwang (MIT) for test specimens and helpful discussions.

REFERENCES

1. R.L. Tobler, R.P. Reed, I.S. Hwang, M.M. Morra, R.G. Ballinger, H.

Nakaj ima, and S. Shimamoto, S., Charpy Impact Tests Near Absolute

Zero, ASTM Journal of Testing and Evaluation, accepted for publication

in Jan, 1991

.

2. Standard Test Methods for Notched Bar Impact Testing of Metallic

Materials, Designation E 23-88, 1989 Annual Book of ASTM Standards,

Vol. 03.01, Amer. Soc. Test. Maters., Philadelphia, 1989, p. 198-213.

3. M.M. Morra, I.S. Hwang, R.G. Ballinger, M.M. Steeves, and M.M. Hoenig,

Effect of Cold Work and Heat Treatment on the 4 K Tensile Fatigue and

Fracture Toughness Properties of Incoloy 908, Proceedings of the

Eleventh International Conference on Magnet Technology, MT-11, Aug 28,

1989, Tsukuba, Japan.

4. S. Jin, W. A. Horwood, J.W. Morris, Jr., and V.F. Zackay, A Simplified

Method for Charpy Impact Testing Below 6 K, in: Adv. Cryo. Eng., Vol.

19, Plenum Press, New York, 1974, pp. 373-378.

5. T. Ogata, K. Hiraga, K. Nagai, and K. Ishikawa, A Simplified Method for

Charpy Impact Testing Near Liquid Helium Temperature, Cryogenics, Vol.

22, 1982, pp. 481-482.

6. L.L. Sparks and R.L. Powell, Low Temperature Thermocouples: KP,

"Normal" Silver, and Copper versus Au-0.02 at% Fe and Au-0.07 at% Fe,

Journal of Research of the National Bureau of Standards A. Physics and

Chemistry, Vol. 76A, No. 3, May-June, 1972, pp. 263-283.

282
COMPARISONS BETWEEN J Ic AND CHARPY ABSORBED ENERGY

AS MEASURES OF CRYOGENIC FRACTURE TOUGHNESS

I.S. Hwang, M.M. Morra, and R.G. Ballinger

Departments of Materials Science and Engineering and Nuclear Engineering

Massachusetts Institute of Technology

Cambridge, Massachusetts 02139, U.S.A.

H. Nakaj ima and S. Shimamoto

Naka Nuclear Fusion Establishment

Japan Atomic Energy Research Institute

Naka-machi, Ibaraki-ken, 311-02, Japan

R.L. Tobler

Materials Reliability Division

National Institute of Standards and Technology

Boulder, Colorado 80303, U.S.A.

The Charpy impact test has been used in cryogenics as a simplified method

to determine the fracture toughness of structural alloys at low temperatures.

A major problem in Charpy testing near absolute zero is that adiabatic heating

during high strain rate deformation increases the fracture temperature [1] . In

this paper, we establish experimental correlations between the critical energy

line integral (J Ic ) and Charpy absorbed energy (Cv ) at both 4 and 77 K for

materials ranging from fully austenitic alloys to ferritic steels. Fracture

mechanics tests are used for J Ic determinations and the liquid cryogen flow

method is used for G^ measurements. At 77 K, a good correlation is found between

J lc and Cv . At 4 K the correlation is poor, indicating that G^ is a poor

indicator of static fracture toughness at this temperature. Interestingly, a good correlation is observed between J Ic at 77 K and C, at 4 K. The results are

explained by the temperature rise of Charpy specimens from 4 K to near 200 K

during impact, as predicted earlier from energy balance calculations.

For annealed austenitic stainless steels between 77 and 4 K, Cv changed

little although J Ic decreased significantly. For ferritic steels, both Cv and

J Ic decreased strongly between 77 and 4 K. For a precipitation-hardened

superalloy, both Cv and J Ic were insensitive to temperature in this range. This

demonstrated wide variability of material response depends on the relative

magnitude of thermal softening. The use of Cv may lead to a serious

overestimation of fracture resistance near absolute zero, especially for alloys

having large thermal-softening effects. Therefore, the Charpy test near absolute

zero is regarded as an invalid measure of the static fracture resistance. A C^-

based acceptance criterion developed for one group of alloys will not be validity

for another group.

283
INTRODUCTION

Both linear elastic (KIc ) and elastic-plastic (J Ic ) fracture mechanics

tests and Charpy impact tests have been used in cryogenics to characterize the

cryogenic fracture behavior of structural alloys. Both methods relate to the

energy absorbed during fracture. Temperatures of interest vary from 4.2 K for

nuclear fusion devices and 20 K for aerospace applications to 77 K and higher for

more general cryogenic applications. Since alloys with plastic fracture behavior

are selected for these applications, the J Ic test is routinely used. The J Ic test

established as ASTM Method E 813-88 can be performed in cryogenic environments

with special modifications to testing machines. Recently, a proposed standard

cryogenic test procedure was developed that employs slow loading to minimize the

crack tip temperature rise [2]. The procedure can be used to determine the

critical energy line integral, J Ic for crack growth, using a compact specimen.

The single specimen technique using unloading compliance for crack size

measurement has been successfully used for a wide variety of structural alloys

including tough austenitic alloys. Often the plane strain fracture toughness is

estimated from

(Kic ) 2 J t E

Ic

1-a2

( 1 )

where E is the elastic modulus and a is Poisson's ratio. For cleavage fractures,

Method E 813-88 is not appropriate; then, the plane strain fracture toughness

test specified by ASTM Method E 399 may be used and J Ic can be estimated from KIc using Eq . (1)

.

The impact test, ASTM Methods E 23-88, was originally developed as a simple

measure of toughness for ferritic steels exhibiting ductile-to-brittle transition

behavior. It is especially meaningful if the Charpy absorbed energy (Cv ) and the

plane strain fracture toughness can be correlated. Charpy tests, require

significantly less complicated apparatus and a much smaller specimen size (10 x 10 x 55 mm) than fracture mechanics tests. However, the Charpy specimen's

temperature will rise during transfer from a helium Dewar to the anvil and during

impact loading, which is a serious drawback. We showed recently [1] that

although the temperature rise during specimen transfer can be minimized by

various cooling techniques, the unavoidable adiabatic heating effect at impact

is too significant to be ignored, especially at start temperatures below 77 K.

The effect of the temperature rise on Cv measurements can be two-fold: at least

large data scatter, and more seriously a lack of correlation with Klc or J Ic .

In this paper we present Charpy impact and fracture mechanics test results

at 77 and 4 K for various structural alloys ranging from ferritic steels to fully

austenitic alloys. Correlations are developed between J Ic and Cv for alloys

representing a broad range of toughness. The results demonstrate the limitations

of Cv as a measure of cryogenic toughness below 77 K.

MATERIALS

Eight structural materials, including 4 fully austenitic alloys, two

metastable austenitic stainless steels, and two ferritic steels, were tested.

The material designations and compositions are listed in Table 1 in the order of

decreasing nickel content.

284
Incoloy 908 [4] was supplied in the form of 27 mm plate in the solution

annealed condition; the material was then aged for 200 h at 650°C in vacuum.

Stainless steel SUS 310 was a commercial heat of 25 mm plate in the solution

annealed condition [5]. JN1 was supplied in the form of 30 mm plate taken from

the surface of a 100 mm hot-rolled plate [3].

AISI 316LN was part of the forged cylindrical vessel of Alcator C-MOD [6]

being constructed at the Plasma Fusion Center of Massachusetts Institute of

Technology (MIT) . The vessel was forged at 1400 K and solution annealed at 1330

K for 14 h and water-quenched. Then a 12% cold work was applied and the alloy

was stress—relieved at 873 K for 20 h and furnace cooled. Material for test

specimens was taken from one of the pin holes drilled in the radial direction of

the machined vessel which had a 190 mm thick wall.

Metastable austenitic stainless steel 304LN, from an earlier experimental

program at the National Institute of Standards and Technology (NIST) [ 7 ]

,

was a

modified heat containing 3.85% Mn; it was produced from a 25 mm plate that was

hot-rolled at 1450 K and solution annealed at 1340 K for h h.

Ferritic 5Ni steel was obtained as a 32 mm plate in the austenitized,

temperized, and reversion annealed condition according to ASTM specification A

645-74 [8] . Fe-18Cr-3Ni-13Mn steel was supplied as 25 mm plate in the hot-rolled

condition. This alloy is of interest because although it is austenitic, it fails

in static tests at 4 K by slip-band cracking which is a form of cleavage. Ultralow carbon steel was supplied in the form of 25 mm plate in the hot-rolled

condition [9]. The tensile and elastic properties for the eight alloys at

selected temperatures are summarized in Table 2.

EXPERIMENTAL PROCEDURES

Tension and fracture toughness properties for some of the alloys were

available from previous tests at NIST; additional tests were performed using a cryogenic facility at MIT. The MIT test facility consists of a servohydraulic

fatigue machine equipped with a specially designed load frame. The loading

system, with a maximum capacity of 90 kN, operates with a commercial Dewar. A

schematic of the cryostat is shown in Fig. 1. Total load train compliance is

estimated to be 2.6 x 10-5 mm/N. The load cell was calibrated to an accuracy of

0.25% of full scale (90 kN) . Specimen strain or crack opening displacements were

measured using ring-shaped extensometers designed for use down to 4 K. The

extensometers were calibrated on a weekly basis to a precision of 2.5 nm using

a dial gage.

Compact specimens and Charpy specimens with notches in the T-L orientation

were machined from the same blocks of materials. Two specimens of each type were

tested. The 25 mm thick compact specimen is a standard design recommended in

Method E 813-88. Most specimens were precracked at room temperature with a final

of 33 MPa m h and stress ratio of 0.1. The ultra-low carbon steel specimens

were precracked with a final aK of 20 MPam* to minimize any warm prestress

effect on the fracture toughness at low temperatures [9]; they were also stress

relieved at 550°C for 1 h in vacuum. The stress relief can remove up to 80% of

residual stress [10].

For all materials the initial crack lengths were 64 to 70% of specimen

width. A clip gage extensometer was mounted at the load-line for measurement of

285
Table 1. Chemical compositions of test materials (mass %).

Material Zft Mi £l £

Incoloy 908 40.6 49.7 3.83 0.01

SUS 310 bal. 19.64 24.58 0.06

JN1 bal. 15.26 25.2 0.018

AISI 316LN bal. 13.10 16.70 0.02

304LN (+Mn) bal. 7.84 19.25 0.024

5Ni steel bal

.

5.03 - 0.08

Fe-18Cr-3Ni-13Mn bal

.

3.26 18.09 0.038

Ultra-low bal. 0.00

carbon steel

A1 Mb 11 MS M Z £1

0.04 1.04 2.99 1.58 - - - 0.14

1.52 - - - - - 0.013 0.37

4.1 - - - - 0.37 0.019 0.33

1.52 - - - 2.58 0.19 0.02 0.41

3.85 - - - 0.31 0.21 0.022 0.65

0.60 0.08 - - 0.30 0.01 0.01 0.25 , 13.32 - - - 0.12 0.37 0.028 0.52

0.021 0.052 - — _ 0.003 0.012 0.001

load cell

Dewar

head plate

pull rod

SIDE VIEW

crosshead

upper reaction

cage rods

3K

lower reaction

cage rods

test specimen

reaction plate

Figure 1 . The MIT cryostat for tension and fracture tests

286
Table 2. Tensile and elastic properties of test materials.

Material Temoerature IS. UTS Elgngfrtjpn Elastic m (MPa) (MPa) ill Modulus (GPa)

Incoloy 908 295 961 1354 13 179

77 182

4 1070 1780 20 182

SUS 310 295 225 558 52 190c 4 796 1322 80 200c

JN1* 295 420 813 52 190c 77 1031 1599 46 200c 4 1336 1812 30 203

AISI 316LN 295 303 664 56 187

77 684 1304 58 174

4 889 1369b 56b 207

304LN (+Mn) 295 329 645 56 190c 76 781 1474 28 200c 4 986 1668 38 200c

5Ni steel 295 534 675 34 198

76 738 1124 22 208

4 1103 1103 “1 209

Fe-18Cr-3Ni-13Mn 295 440 796 56 195

76 1144 1519 18 196

4 1540 1811 4 196

Ultra-low 295 137 268 49 207

carbon steel 76 758 793 1.6 211

4 705 744 1.2 211

* Data are measured at 4 K and estimated from Ref. [3] at 295 and 77 K.

b Liquid helium level was low during plastic deformation.

c Estimated values.

287
displacement. Crack length was determined from the unloading compliance using

the equations of Saxena and Hudak [11]. Specimen rotation and moving crack

effects were taken into account in the crack length determinations. Tests were

performed with a fully automated procedure at a constant gage displacement rate,

typically 0.1 mm/min. All testing was in accordance with Method E 813-88. Data

analysis was in accordance with ASTM 1152—87. The J Ic test system was verified

by interlaboratory tests using two austenitic stainless steels with known

fracture toughness [12].

To complete the database, tension tests were performed where necessary

using specimens 6.35 mm in diameter and 25.4 mm in gage length. Tests were

according to the proposed ASTM standard for tension tests at liquid helium

temperature. The specimens were axially oriented in the same direction as the

loading axis of the fracture toughness specimens. All tests were straincontrolled at 2 x 10-4

s* 1 by using an extensometer with 25.4 mm gage length.

The Charpy test specimens were full size (10 x 10 x 55 mm) with a 2 mm deep

notch in strict accordance with ASTM Method E 23-88. Most tests were performed

with a standard C-type pendulum machine at the Superconducting Magnet Laboratory

of the Japan Atomic Energy Research Institute (JAERI) . Some tests at 4 K used

a C-type pendulum machine at NIST, and some at 77 K used a U-shaped pendulum

machine at MIT. There was an overlap in the testing program at the three

institutions so that a cross check of data can be made. The impact test machines

meet the requirements of Method E 23-88 with a full scale of 300 J, except for

the MIT machine which was not firmly attached to the floor. This influenced the

MIT data slightly, as discussed later.

The 77 K Charpy tests at JAERI and MIT followed the standard Method E 23-88

procedure: conventional cooling of specimens in a Dewar for about 10 min, and

subsequent transfer to the anvil within 5s. As pointed out previously [1],

there is no standard test procedure for Charpy testing at 4 K. Therefore, two

test techniques at 4 K were considered: the glass Dewar method [1,13] and the

flow method [1,14]. The advantages and disadvantages of each technique were

studied and will be reported elsewhere in detail. In this paper, Charpy data

from the flow method were preferred because that method applies equally well to

both low and high toughness materials, and no calibration correction of Cv data

is required. In the flow method, the specimen is loosely wrapped with

polyethylene sheet of 0.5 mm thickness. The specimen is secured in the wrap

using double-sided adhesive tape; then the assembly is connected to the cryogenic

transfer line. The specimen in a continuous flow of helium cools to within 2 K

of the intended start temperature in approximately 2 min. The specimen cooling

time and start temperature (5 K ± 1 K) were monitored and confirmed by

calibrations using a gold-iron thermocouple. In this paper we call absorbed

energy measurements made by this method "Cv at 4 K."

RESULTS AND DISCUSSION

The J Ic and Cv measurements at 77 and 4 K are listed in Table 3. Some J Ic values were estimated from KIc data reported earlier. SUS 310 stainless steel

was tested at 4 K only due to shortage of stock material. Type 316 LN stainless

steel in J Ic tests at 77 K did not exhibit the required crack growth and a valid

J Ic was not obtained at this temperature. The 4 K Charpy test results at JAERI

and NIST agree within 5% as shown in Table 3. The 77 K Charpy test results at

JAERI and MIT agree within 15% . The larger deviation in the latter case is

288
Table 3. Results of fracture toughness and Charpy impact tests.

Jr. (kJ/m2 Charov Absorbed Enerzv (J)

Material ZZ_£ ii 77 K

Incoloy 908 243 ± 6 235 ± 5 63 ± 2(a) 70 ± 1

SUS 310 - 391(c) — 166 ± 2

JN1 600 225 ± 22 274(b) 235 ± 3

AISI 316LN (d) 576 ± 22 255 ± 5

279 ± 5(b)

231 ± 7

231 ± 9(e)

304LN (+Mn) 332 ± 7 101(c) 102 ± 6 94 ± 3

99 ± 3

5Ni steel 93 ± 1(f) 15.2 ± 2 . 7 (c ) 39 ± 4(b) 17 ± 1

Fe-18Cr-3Ni-13Mn 227 ± 2 31.2(c) 80

69 ± 4(a)

80 ± 4(b)

45 ± 2

Ultra-low 3.3 ± 0.5(c) 0.9(f) 2.7(b) 8 ± 1

carbon steel

Note

(a) Method E 23-88 tests conducted at JAERI

.

(b) Method E 23-88 tests conducted at MIT.

(c) Estimated from KIc using Eq. (1).

(d) Required crack extension for valid test was not obtained.

(e) Flow method tests conducted at NIST.

289
significant for high C^,; presumably, it is due to less firm attachment of the

testing machine to the floor at MIT. Although the difference is small, JAERI's

data are used In the correlations with J Ic .

The J Ic - Cy correlations at 77 and 4 K are presented in Figs. 2 and 3,

respectively. Each figure shows a solid parabolic trend line drawn by least

squares fitting. At 77 K,

Cv - 10.4 + 0.136(J Ic ) + 0 . 04996(J Ic ) 2 with r

2 - 0.988, (2)

and at 4 K,

Cv - 18.0 + 0 . 6708 (

J

Ic ) - 0 . 05562 (J Ic ) 2 with r

2 - 0.727, (3)

| where r

2 is the correlation coefficient. Based on the correlation coefficients

and data scatter shown in our figures, the correlation between J Ic and at 77

K is good, but the correlation at 4 K is poor.

The poor correlation at 4 K cannot be attributed to property measurement

errors (they change little between 77 and 4 K) , but is attributable to large

temperature rises in the Charpy test specimens. Below 77 K, the specific heats

of structural alloys are extremely low. Because the specimens heat during high

strain rate deformation, we expect fracture temperatures much higher than the

start temperature. As shown in Fig. 4 from previous work [1], temperature

increases as high as 180 K are predicted and measured for Charpy specimens

initially at 4 K. The predicted maximum temperature rises for alloys are nearly

the same for different starting temperatures ranging from 77 to 4 K.

Assuming adiabatic heating is the main reason for the poor correlation

between Cv and J Ic at 4 K, and that the behavior predicted in Fig. 4 is accurate,

it follows that:

(1) Impact tests at 4 K actually involve fractures near 180 K, and a fair correlation is expected between Cv at 4 K and J Ic at 77 K.

(2) Regardless of the Charpy test start temperature (77 or 4 K) , the

maximum specimen temperature reached during fracture and the

measured values will be similar for a given alloy, even if J Ic at

these temperatures is quite different.

(3) For alloys exhibiting transitional fracture behavior, scatter should

increase if there is a difference in failure mechanisms between

I Charpy and compact specimens. Charpy specimens of certain alloys

may exhibit ductile tearing rather than cleavage if they heat during

impact to temperatures above the ductile—to—brittle transition

temperature

.

The first hypothesis is examined by Fig. 5 where C, at 4 K is plotted

versus J Ic at 77 K. A least-squares fit gives

(Cv )4 K - 5.43 + 0. 1158(J Ic )77 K + 0 . 000446 [ (J Ic ) 77K] 2

, (4)

with r

2 - 0.994.

290
Figure 2. Relationship between J and Charpy absorbed energy at 77 K.

Figure 3. Relationship between J and Charpy absorbed energy at 4 K.

Ic

291
Final

Temperature,

Figure 4. Predicted maximum temperature rise for Charpy specimens

as a function of initial temperature (from Ref. [1]).

292
Figure 5. Relationship between J at 77 K and Charpy absorbed

energy at 4 K.

8

6

4

2

0

0 2 4 6 8

(C v at 77 K)/(C v at 4 K)

Figure 6. Comparison of and after normalization to 4 K values.

293
This correlation is much better than that for and J Ic at 4 K given in

Eq. (3). The improvement in the correlation from Fig. 3 and Eq. (3) to Fig. 5

and Eq. (4) suggests that the 4 K Charpy specimens fractured at a higher

temperature than JIc specimens.

The correlations in Figs. 2 and 5 are quite similar because there is little

difference in Cv at 77 and 4 K. To show this point clearly, in Fig. 6 the ratio

of Cv at 77 K to that at 4 K is plotted vs. the ratio of J Ic at 77 K to that at

4 K. For alloys at 77 and 4 K, Cv does not change by more than a factor of 2,

whereas J Ic increases by factors up to 7. The results support our second

hypothesis

.

Two steels that cleave at 4 K and show large changes in J Ic between 77 and

4 K were examined under the scanning electron microscope. The ferritic 5Ni and

austenitic Fe-18Cr-3Ni-13Mn steels show changes in fracture morphology depending

on the type of test performed. For 5Ni steel, Figs. 7 and 8 show representative

fracture surfaces of Charpy and compact specimens tested at 77 and 4 K. Cleavage

predominates in both the Charpy impact and fracture toughness test specimens at

77 and 4 K, as shown in Fig. 7. Unlike the fracture toughness specimens,

however, the Charpy specimens include narrow regions at the notch tip where

microvoid coalescence occurs. Figures 9 and 10 show fracture surfaces for Fe18Cr-3Ni-13Mn stainless steel. In this case, the fractographic features of

Charpy specimens tested at high strain rates are quite unlike those for compact

specimens at low strain rates. Figure 9 shows that the Charpy specimen fracture

surfaces are entirely microvoid coalescence type ductile tearing at both test

temperatures. In contrast, Fig. 10 shows that the J Ic specimens exhibit a

substantial portion of cleavage, although the percentage is higher at 4 K than

at 77 K at 77 K. Fibrous fracture also occurs as an intermediate fracture mode

between cleavage and microvoid coalescence modes. Thus, fractographic evidence

for Fe-18Cr-3Ni-13Mn indicates that the Charpy specimens, initially at 4 K,

heated above 77 K to temperatures at which the fracture mechanism is microvoid

coalescence. The point is that a good correlation between J Ic and cannot be

expected if different test specimens of the same alloy fail by different

mechanisms

.

For annealed austenitic stainless steels JN1 and AISI 304 LN between 77 and

4 K, J lc changed by factors of 2 to 3, while Cv changed little as shown in Table

3. In contrast, the aged superalloy shows almost no change in J Ic or between

77 and 4 K. This temperature insensitivity is expected because the tensile

properties of the alloy are also temperature insensitive [14], Indeed, the

mechanical properties of precipitation-hardened alloys are generally less

temperature dependent than solution treated or annealed alloys, and the relative

magnitude of the thermal softening increases from the precipitation-hardened

superalloys, to austenitic stainless steels, and finally to ferritic steels.

This variation in thermal-softening among the studied alloys is shown in Fig. 6

by the ratio of J Ic measurements at two temperatures. Variations in thermal

softening of alloys is the primary reason why a general correlation between J Ic and Cv below 77 K can not be obtained.

It appears that Charpy tests are of little value for characterizing the

static fracture toughness at 4 K. At best, we might attempt to use the 4 K

Charpy test for screening of brittle materials, as in regulatory protocols. The

attempt can be seriously faulted due to the wide variation in the thermal

softening of alloys. If an acceptance criterion were to be derived from data for

294
Figure 7. SEM fractographs of Charpy specimens: 5Ni steel tested at 4 K

(upper, 500X) and at 77 K (lower, 500X).

295
Figure 8. SEM fractographs of fracture toughness specimens: 5Ni steel

tested at 4 K (upper, 500X) and at 77 K (lower, 500X)

.

296
Figure 9 . SEM fractographs of Charpy specimens: Fe-1 8Cr-3Ni-1 3Mn steel

tested at 4 K (upper, 500X) and at 77 K (lower, 500X).

297
Figure 10. SEM fractographs of fracture toughness specimens of

Fe-18Cr-3Ni-1 3Mn steel specimens: tested at 4 K (upper, 500X)

and at 77 K (lower, 500X)

.

298
a particular group of alloys , it would be quite inappropriate for screening other

alloys. Depending on circumstances, the criterion would be either

overconservative or underconservative . Suppose that a 4 K (^-based acceptance

criterion were to be established for annealed austenitic stainless steels in

accordance with a J Ic-based acceptance criterion. Since annealed austenitic

stainless steels have significant thermal softening, the same criterion would be

unduly restrictive for precipitation-hardened steels or nickel-based superalloys

that have negligible thermal softening.

In using G^ as a measure of energy absorption at extreme cryogenic

temperatures, a dangerous overestimation is possible due to the thermal

softening. Certain alloys in Charpy tests at 4 K may fail by ductile fracture

mechanisms even though the normal fracture mechanism at slow loading rates below

77 K is cleavage. The risk of overestimation, the poor correlation with J lc , and

the lack of quantitative data for use in design makes the Charpy impact test

unacceptable as a measure of fracture toughness near absolute zero.

CONCLUSIONS

(1) Charpy test specimens with an initial temperature at or near 4 K actually

fracture at temperatures well above 77 K due to adiabatic heating. We

observe the following consequences:

(a) There is a poor correlation between G* and J Ic at 4 K, and a good

correlation between G* at 4 K and J Ic at 77 K.

(b) There is little difference in G* measurements for alloys at 77 and

4 K, even if J lc changes by several factors.

(c) For some alloys such as Fe-18Cr-3Ni-13Mn stainless steel, Charpy

specimens tested at an initial temperature near 4 K show different

fractographic features than slowly loaded compact specimens tested

at 4 K.

(2) High strain rate tests including the Charpy impact test are of limited

value for predicting the static toughness below 77 K. Even for screening,

an acceptance criterion for one group of alloys would significantly

overestimate or underestimate the fracture toughness for another group due

to wide variations in thermal softening among alloys

.

REFERENCES

[1] Tobler, R.L.

,

Reed, R.P., Hwang, I.S., Morra, M.M.

,

Ballinger, R.G.,

Nakajima, H. , and Shimamoto, S., "Charpy Impact Tests Near Absolute Zero,"

Journal of Testing and Evaluation, in press.

[2] Tobler, R.L., "A3. Test Procedures for Austenitic Alloys in Liquid Helium

at 4 K," annex proposed for ASTM method E 813-88, presented to

subcommittee E 24.04 at ASTM Technical Committee Week Meeting, San

Francisco, April 24, 1990.

[3] Nakajima, H. , Yoshida, K. , Shimamoto, S., Tobler, R.L.

,

and Reed, R.P.,

"Round Robin Tensile and Fracture Toughness Test Results for CSUS-JNl(Fe299
25Cr-15Nl-0. 35N) Austenitic Stainless Steel at 4 K," in Adv. Cryo. Eng.,

Vol. to be published.

[4] Morra, M.M. , M.S. Thesis, submitted to Massachusetts Institute of

Technology, 1990.

[5] Tobler, R.L., Shoji, T. , Takahashi, H. , and Ohnishi, K. , "Fracture,

Acoustic Emission and Adiabatic Heating of Austenitic Stainless Steels at

Liquid Helium Temperature," in Acoustic Emission III, The Japanese Society

of NDI, Tokyo, 1986, pp. 453-461.

[6] Ballinger, R.G.

,

Becker, H.D.

,

Hwang, I.S., and Morra, M.M. , "Superstructure for Super Tokamak," 13th IEEE Symposium on Fusion

Engineering, Knoxville, TN, 1989.

[7] Tobler, R.L., and Reed, R.P., "Tensile and Fracture Properties of

Manganese-Modified AISI 304 Type Stainless Steel," in Adv. Cryo. Eng.,

Vol. 28, Plenum, New York, 1982, pp. 83-92.

[8] Tobler, R.L.

,

Mikesell, R.P., Durcholz, R.L.

,

and Reed, R.P., "Low

Temperature Fracture Behavior of Iron-Nickel Alloy Steels" in Properties

of Materials for LNG Tankage, ASTM STP 579, American Society for Testing

and Materials, Philadelphia, 1975, pp. 261-287.

[9] Tobler, R.L., Ma, L.M.

,

and Reed, R.P., "Mechanical Properties and Warm

Prestress of Ultra-Low Carbon Steel at 4 K," presented in 12th Cryogenic

Structural Materials Workshop, Vail, CO, October 11-12, 1989.

[10] Salkin, R.V.

,

"The Stress Relaxation Heat Treatments after Welding," Proc.

Int. Conf. on Stress Relieving Heat Treatments of Welded Steel

Constructions, Sofia, Bulgaria, July 6-7, 1987.

[11] Saxena, A., and Hudak, S.J., Jr., "Review and Extension of Compliance

Information for Common Crack Growth Specimens," Int. J. Fracture, Vol. 14,

No. 5, 1978, pp. 453-468.

[12] Morra, M.M.

,

Hwang, I.S., Ballinger, R.G., Steeves, M.M.

,

Hoenig, M.O.,

"Effect of Cold Work and Heat Treatment on the 4 K Tensile, Fatigue, and

Fracture Toughness Properties of Incoloy 908," presented at 11th

Interntional Conference on Magnet Technology, August 28-September 1, 1989.

[13] Mori, K. , and Kuroda, T. , "Prediction of Energy Absorbed in Impact for

Austenitic Weld Metals at 4 K," Cryogenics, Vol. 25, 1985, pp. 243-248.

[14] Ogata, T. , Nagai, K. , Hiraga, K. , Ishikawa, K. , and Irie, H. , "Effect of

EB-Weld and Cold Rolling on Low Temperature Strength and Toughness of

Austenitic Stainless Steels," in Austenitic Steels at Low Temperatures,"

R.P. Reed and T. Horiuchi, Eds., Plenum, New York, 1983, pp. 211-220.

300
JAPAN—U.S . COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

NINTH DRAFT

NOVEMBER 1989

NOTE: This is a working document of a standard being developed by

ASTM subcommittee E 28.10; it is intended for committee review,

but not for general distribution at this time.

Proposed Standard Test Method for

TENSION TESTING OF STRUCTURAL ALLOYS IN LIQUID HELIUM

1 . Scope

1.1 This standard describes procedures for the tension testing of

structural alloys immersed in liquid helium. The format is similar to that

of other ASTM tension test standards, but the contents include modifications

for cryogenic testing, which requires special apparatus, smaller specimens,

and concern for serrated yielding, adiabatic heating, and strain- rate

effects

.

1.2 To conduct a tension test by this standard, the specimen in a cryostat is fully submerged in normal liquid helium (He I) and tested using

crosshead displacement control at a nominal strain rate of 10“ 3 s _1 or less.

Tests using load control or high strain rates are not considered.

1.3 This standard specifies methods for the measurement of yield

strength, tensile strength, elongation, and reduction of area. The determination of the elastic modulus is treated in Method E 111.

1.4 Values stated in SI units are treated as primary; values stated in

U.S. customary units are treated as secondary.

1.5 This standard may involve hazardous materials, operations, and

equipment. This standard does not purport to address all of the safety

problems associated with its use. It is the responsibility of the user of

this standard to establish the appropriate safety and health practices and

determine the applicability of regulatory limitations before use.

NOTE 1 —The boiling point of normal liquid helium (He I)

at sea level is 4.2 K (-452. 1°F or 7.6°R). It decreases with

geographic elevation and is 4.0 K (-452. 5°F or 7.2°R) at the

National Institute of Standards and Technology in Colorado,

1677 m (5500 ft) above sea level. In this standard the

temperature is designated 4 K.

301
2. Applicable Documents

2 . 1 ASTM Standards :

A 370 Methods and Definitions for Mechanical Testing of Steel Products

( Annual Book of ASTM Standards , Voi. 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 74 Standard Practice of Calibration of Force -Measuring Instruments

for Verifying the Load Indication of Testing Machines

(Annual Book of ASTM Standards , Vol. 03.01)

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)

E 1012 Standard Practice for Verification of Specimen Alignment Under

Tensile Loading (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

302
Z 8103 G lossary 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.

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 tensile forces to

test specimens in cryogenic environments. A schematic illustration is shown

in Fig. 1

.

LOAD

t

FIG. 1 Typical Cryostat for Tension Testing at 4 K.

303
3.1.2 Dewar—a vacuum- insulated container for cryogenic fluids.

3.1.3 adiabatic heating—the internal heating of a specimen resulting

from tension testing under conditions such that the heat generated by plastic

work cannot be quickly dissipated to the surrounding cryogen.

3.1.4 reduced section—section in the central portion of the specimen,

which has a cross section smaller than the gripped ends.

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 longitudinal strains measured at

opposite or equally spaced surface locations on the sides of the longitudinal

axis of symmetry of the specimen. The longitudinal strains are measured

using two or more strain- sensing devices located at the mid- length of the

reduced section.

3.1.9 bending strain—the difference between the strain at the surface

of the specimen and the axial strain. The bending strain varies around the

circumference 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 two, three, or more circumferential positions, and at each of two

different longitudinal positions.

3.1.11 discontinuous yielding stress, a—the stress at the initiation of

the first measurable serration on the curve of stress -versus -strain. The

parameter a is a function of test variables and is not a material constant.

4. Significance and Use

4.1 Tension 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 The load—time and load—deflection records for alloys tested in

liquid helium using displacement control are serrated (1). Serrations are

formed by repeated bursts of unstable plastic flow and arrests. The unstable

plastic flow (discontinuous yielding) is a free-running process occurring in
localized regions of the reduced section at higher than nominal rates of

strain vith internal specimen heating. Examples of serrated stress-strain

curves for a typical austenitic stainless steel with discontinuous yielding

are shown in Fig. 2.

4.3 A constant specimen temperature cannot be maintained at all times

during tests in liquid helium. The specimen temperature at local regions in

the reduced section rises temporarily above 4 K during each discontinuous

yielding event (see Fig. 2), owing to adiabatic heat. The number of events

and the magnitude of the associated load drops are a function of the material

composition and other factors such as specimen size and test speed.

Typically, altering the mechanical test variables can modify but not

eliminate the discontinuous yielding (2-4). Therefore, tensile property

measurements of alloys in liquid helium (especially tensile strength,

elongation, and reduction of area) lack the usual significance of property

measurements at room temperature where deformation is more nearly isothermal

and discontinuous yielding typically does not occur.

FIG. 2 Typical Engineering Stress—Strain Curves and Specimen

Temperature Histories, at Four Different Nominal Strain

Rates, for AISI 304L Stainless Steel Tested in Liquid

Helium (Ref. 4).

305
4.4 The stress-strain response of a material tested in liquid helium

depends on whether load control or displacement control is used (3).

Crosshead displacement control is specified in this standard since the goal

is material characterization by conventional methods. The possibility of a

different and less favorable material response must be taken into account

when data are used for design in actual applications subject to loadcontrolled conditions.

5 . Hazards

Several precautions must be observed in the use of cryogenic fluids and

equipment. Skin or eye contact with cryogens will produce burns. The

appropriate protection may require goggles, clothing without pockets or

cuffs, gloves, and tongs for handling cold specimens. Cryogenic containers

that are internally pressurized or evacuated are potentially hazardous in

that damage or leaks can produce explosions or implosions. Also, when

liquids evaporate to gases, there is a huge volume increase; therefore

asphyxiation is a potential threat where liquid nitrogen or liquid helium

evaporates in rooms that are not properly ventilated. Safety guidelines

pertaining to the use of liquid helium and other cryogenic fluids are

considered elsewhere in more detail (5)

.

6 . Apparatus

6.1 Test Machines—Use a test machine that meets the requirements of

Practices E 4 regarding verification of force accuracy. The test machine

compliance (displacement per unit of applied force of the apparatus itself)

should be known. The compliance can be measured by coupling the load train

without including a specimen, by replacing the specimen with a rigid block,

or by using a special calibration specimen. The compliance is then measured

at a low force and at the highest force used to qualify the machine, as

directed in section 6.4.1 of this standard.

6.2 Syscem Design—Typically, alloys in liquid helium exhibit double or

triple their ambient strengths. For the same specimen geometry, higher

forces must be applied to the cryostat, test specimen, load train members,

and grips at cryogenic temperatures. Since many conventional test machines

have a maximum force of 100 kN (22 480 lbf) or less, it is recommended that

the apparatus be designed to accommodate one of the small specimens cited in

section 8.2.2 of this standard.

6.3 Construction Materials—Many construction materials, including the

vast majority of ferritic steels, are brittle at 4 K. To resist failure in

service, fabricate the grips and other load-train members using strong,

tough, cryogenic alloys. Materials that have low thermal conductivity are

desirable to reduce heat flow. Austenitic stainless steels (AISI 304LN)

,

maraging steels (200, 250, or 300 grades, with nickel plating to prevent

rust), and extra- low- interstitial (ELI) grade titanium alloys (Ti-6Al-4V and

Ti-5Al—2 . 5Sn) have been used with proper design, for grips, pull rods, and

306
cryostat frames. Nonmetallic materials (for example, glass—epoxy composites)

are excellent insulators and are sometimes used for compression members.

6.4 Alignment :

6.4.1 Single- Specimen Apparatus—Proper system alignment is essential to

avoid bending strains in tension tests. For a conventional single- specimen

cryostat, the machine and grips should be capable of applying force to a

precisely machined calibration specimen so that the maximum bending strain

does not exceed 10% of the axial strain. Reduce bending strain to an

acceptable level by making proportional adjustments to a cryostat having

alignment capability, or by using spacing shims to compensate an unadjustable

fixture. Calculation the strain based on readings taken while the

calibration specimen is subjected to a low force, as well as at the highest

force for which the machine and load train are being qualified. Procedures

for measuring specimen alignment are given in Standard Practice E 1012.

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 amount of bending

strain owing to small variations in the proposed test specimen

configurations, or differences in machining.

6.4.2 Multiple-Specimen Apparatus—For this type of cryostat the

alignment depends on the type of fixtures used. Measure and reort the

maximum bending strain.

6.4.3 Qualify the apparatus by making 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, and the

specimen concentricity should be as nearly perfect as possible. To the limit

of strength for the alignment specimen, 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.

6. 4. 3.1 For cylindrical specimens, calculate the maximum bending strain

defined in 3.1.10 from the strains measured at three circumferential

positions, at each of two different longitudinal positions. Measure the

strains with three electrical-resistance strain gages, extensometers , or clip

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.

6.4. 3.2 For specimens of square or rectangular cross section, measure

strain at the center of each of the four sides or, in the case of thin cross

sections, near the outer edges of each of the two broad sides.

6. 4. 3.

3

To eliminate the effect of specimen bias, repeat the axiality

measurements with the specimen rotated 180°, but with the grips and pull rods

retained in their original positions. Then calculate the maximum bending

307
strain and the strain at the specimen axis as the average of the two readings

at the same position relative to the machine.

6.4.4 Strain -Averaging Technique—Nonaxiality of loading (which may be

introduced due to the machining of the test specimens) is usually sufficient

to introduce errors in tension tests at small strains when strain is measured

at only one position on the specimen. Therefore measure strains at three

equally spaced (or, if good alignment has been achieved, at least two

opposing) positions within the reduced section. Report the average of the

strains from the two or three positions centered on the reduced section.

6.5 Gripping Mechanisms—The choice of gripping mechanism to be used

is influenced by specimen type. The mechanisms described in Methods E 8 and

E 8M are satisfactory at 4 K, but cryogenic materials must be used in the

construction of components to avoid failure in service.

6.6 Dimension-Measuring Devices—For measuring the dimensions of

specimens, use a micrometer or other device that is accurate and precise to

at least one-half of the smallest unit to which a given dimension must be

measured.

6.7 Cryostats and Support Apparatus

6.7.1 Cryostats—A cryostat capable of retaining liquid helium is

required. 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 (6-10).

6.7.2 Dewars—Stainless steel Dewars are safer (that is, more fracture

resistant) than glass Dewars and less expensive than fiberglass Dewars.

Generally, a single helium Dewar (see Fig. 1) is sufficient for short-term

tensile tests. Also possible is a double-Dewar arrangement in which an

outer Dewar of liquid nitrogen surrounds the inner Dewar of liquid helium.

6.7.3 Ancillary Equipment—Dewars and transfer lines for liquid helium

must be vacuum insulated. Vacuum pumps, pressurized gas, and liquid nitrogen facilities are therefore required. After testing, the helium may be

released to the atmosphere, recycled as a gas, or reliquefied. Recycling

or reliquefaction requires large investments in purification and support

systems

.

6.8 Temperature Maintenance and Liquid-Level Indicators—The intended

test condition is ensured by maintaining a liquid helium environment. With

the specimen completely immersed, a thermocouple to measure its temperature

is not required for routine tests. Instead, a simple indicator or meter is

required to ensure that the specimen remains fully submerged throughout the

test. An on—off indicator of the carbon- resistor type located at some

reference point in the cryostat may be used to verify that the liquid level

always remains above the specimen. Alternatively, the liquid level may be
continuously monitored using a superconducting wire sensor of appropriate

length positioned vertically inside the cryostat.

NOTE 4—One indication of the system nearing and reaching

a steady state condition is the amount of condensation flare.

As liquid helium is transferred into the cryostat, the flare

becomes visible when boiled-off helium contacts room temperature

air at the vent of the cryostat (Fig. 1). As cool-down proceeds,

the flare decreases to a slowly issuing cloud due to less active

boiling as the internal temperature of the cryostat reaches

operating temperature.

6.9 S Crain Gages:

6.9.1 Precautions—Strain-gage films bonded directly to the specimen

surface may be used to measure strain at 4 K (11). The use of bonded strain

gages at 4 K, however, requires precautions not customarily required at room

temperature. There are two major complications: the gage factor varies with

temperature, and thermal output (apparent strain) is introduced as the

specimen—gage combination is cooled from room temperature to 4 K. Thermal

output is caused by two concurrent and algebraically additive effects in the

strain gage installation: (1) the electrical resistivity of the gage grid

element and (2) the differential thermal expansion between the gage grid

element and the test specimen to which the gage is bonded. These effects

must be accounted for, or considerable error in strain measurements may be

introduced.

Gage manufacturers generally do not supply thermal output data at 4 K;

neither do they state gage factors at 4 K. For high accuracy the user may

need to perform gage factor and thermal output calibrations for his system

to establish a stable reference gage output at 4 K before beginning tension

tests. For this reason, strain gage calibrations may be more difficult than

extensometer calibrations (see Section 6.10.3).

6.9.2 Selection and Characteristics—Not every type of strain gage is

usable at cryogenic temperatures. Select a satisfactory combination of gage

active element, backing material, and bonding agent based on experience and

manufacturer's recommendations. A common choice for extreme cryogenic

service is a Ni-Cr-Al-Fe alloy gage with a temperature -compensated active

element (8). Typically the gage resistance is 120 or 350 fl, and a low

excitation voltage is used to reduce Joule heating at 4 K. The full-scale

operating range is typically 1% at room temperature and 2% at 4 K.

6.9.3 Wiring—Various circuits may be used for wiring strain gages.

The choice depends on purpose and accuracy desired. One circuit that is

satisfactory for tension tests according to this standard is the three -wire

temperature -compensating hookup and Wheatstone bridge arrangement shown, for

example, in Fig. 3. The three-wire circuit nullifies or eliminates

thermally induced resistance changes in the leadwires if the wires Rjfo and

R^ 3 in Fig. 3 have the same resistance and experience the same temperature

309
Input

Voltage

FIG. 3 A Method of Strain-Gage Wiring for Cryogenic Tests

changes along their lengths. This circuit may be used with commercial

strain indicators. First, balance the Wheatstone bridge at room

temperature. Then, to compensate for any apparent strain induced on cooling

the specimen and gage to the test temperature, cancel the change in

electrical signal after cooling before the force is applied and testing

begins. Other circuits and instruments are possible, and some techniques

may offer higher accuracy than the illustrated example.

NOTE 5—Some gage manufacturers provide estimated values

of the gage factors for the use of their products at low

temperatures. Their estimated values do not necessarily

agree with published research, and calibration by controlled

experimental determinations are preferred. Gage factors at

temperatures as low as 4 K for some common materials have been

published in a few studies. For example, findings for Ni-Cr

alloy gages show that the gage factor increases nonlinearly by

2.5 or 5% as the temperature is reduced from 295 to 4 K (12- 14).

6.10 Clip-Gage Extensometers

6.10.1 Types—Reliable clip-gage extensometers for use at 4 K may be

purchased or built. An example is the beam gage, which uses four straingage films bonded in a full Wheatstone bridge. Extension within the

specimen gage length is sensed by the detachable extensometer , which is

clipped to retaining pins that are fixed to the specimen reduced section.

6.10.2 Characteristics—To measure the 0.2% offset yield strength, one

or more Class B—2 or better extensometers, as identified in Method E 83, may

be used. Each extensometer must meet the sensitivity and accuracy

requirements of Method E 83 and must be tested to ensure accuracy at 4 K.

Whenever possible, mount the extensometer(s) directly to the specimen

reduced section.

NOTE 6—It may be desirable to use several extensometers

to detect the amount of bending, especially for brittle

310
specimens. One method for obtaining multi-planar strains is

to fasten a pair of collars to the specimen reduced section,

using sharpened radial thumbscrews. The collars are made with

detents (located on the inboard faces) to accept spring-loaded

clip gages. This arrangement also serves to define a fixed

gage length within the specimen reduced cross section. 6.10.3

Calibration-—Calibrate extensometers at room temperature and at

4 K. For calibrations at 4 K, a device such as a micrometer with vertical

extension tubes can be used with the extensometer (s) mounted at the lower

end and immersed in liquid helium. If the calibration is known and proved

to be accurate, linear, and reproducible, then room- temperature checks may

be performed before each series of tests to indirectly verify the 4-K

calibration. However, direct calibration at 4 K must be performed

periodically, especially if damage is suspected or repairs have been made.

6.11 Capacitance Extensometers—Extensometers that use capacitance

measurement to monitor strain may be used (11) . 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 its output is nonlinear, and it must be

compensated for the hyperbolic dependence of capacitance on displacement.

7. Sampling

7.1 Take samples for tension testing from the material in its final

condition to ensure that the properties measured are representative of the

product. Allow for any superficial effects introduced by the cutting

operations

.

7.2 Cut specimens from locations thought to be most representative of

the stock material, realizing that data for specimens taken from selected

locations of a structure or material may not be representative of the whole.

The conventional locations should normally be used:

7.2.1 For products 40 mm (1.6 in) or less in thickness or diameter,

the location should be at the center.

7.2.2 For products over 40 mm (1.6 in) in thickness or diameter, the

location should be midway from the surface to the center.

7.3 Choose a specimen size and shape based on the requirements

necessary to obtain representative samples of the material, and on the test

machine's force capacity (section 6.2).

7.4 Using the notation in section 4 of Method A 370, specify the

orientation of the specimen's axis relative to the most predominant of

either the grain flow or the principal working direction of the final form

of the stock.

311
8 . Test Specimens

8 . 1 General :

8.1.1 Types and Specifications—Any specimen configuration cited in

Methods E 8 or E 8M may be used. Specifications for dimensions, tolerances, and surface finish are stated in sections 6.1 through 6.17

of those standards.

8.1.2 Size—Specimens from sheet or wire products having relatively

small cross-sectional areas can be tested within the force capacities of

conventional apparatus. Specimens from thick plate or bar products,

however, must be machined to a reduced cross-sectional area so the force

capacity of the machine is not exceeded. Preferably, the specimen in

this case should have the same configuration as the standard specimen but

with an appropriately reduced cross section.

8 . 2 Round Bar Specimens :

8.2.1 Standard Room-Temperature Specimens—A 12.5 mm (0.5 in) diameter round bar is the standard configuration for room- temperature tests

according to Methods E 8 and E 8M. Specimens of that size, however,

require high test forces to fracture strong alloys at 4 K. For example,

210 kN (47 208 lbf) is required to test typical AISI 304LN steel at 4 K,

whereas 100 kN (22 480 lbf) is the limit for most test machines.

9 8.2.2 Standard 4 -K Specimens—To meet the load limitations of conventional test machines, the round bar specimens in sections 7. 2. 2.1

and 7. 2. 2.

2

are defined 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 6.4.1 can be met by precise machining.

8.2.2.

1

Standard , small metric specimens

.

These specimens have a

7-mm (0.275-in) diameter and a G-to-D ratio of 5:1.

8. 2. 2.

2

Standard , small U.S. customary specimens

.

These specimens

have a 6.25-mm (0.250-in) diameter and a G-to-D ratio of 4:1.

8.2.3 Alternatives—If the 4-K standard specimens described above are

inappropriate for some reason, 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

Method and Fig. 8 of Methods E 8 or E 8M)

.

8.2.4 Subsize specimens—Special care in fabrication and testing

is required for specimens with diameters less than 6 mm (0.236 in). As the

specimen size is reduced, factors such as machining, surface finish, and

alignment show increasing importance. Also, to ensure polycrystalline

deformation, some experimenters insist on the need for a minimum of ten

grains per cross section; if the number of grains per cross section is

approximately ten or less, note this in the report.

312
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 S tandard

G/D ratio - 5 Small Specimen

(dimensions, mm)

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 Small Specimen

(dimensions, in)

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

1 See also the notes to Fig. 8 in Methods E 8 and E 8M.

FIG. 4 Round Bar Specimen Configuration (see Table 1)

313
99.1. Procedure*

Marking and Measuring the Test Specimen:

9.1.1 Gage Length—Gage marks may be lightly punched, scribed, or inked

at appropriate locations on the reduced section of the specimen. This is

the preferred method. After marking the gage length, measure it to the

nearest 0.05 mm (0.002 in). The conventional gage length is five times the

diameter for metric specimens or four times the diameter for U.S. customary

specimens. If another gage length is used for elongation measurements,

describe it in the report.

NOTE 7—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, coat

the reduced section with layout ink, and mark the gage length

by rotating the specimen in a jig with knife edges scraping off

the ink at the appropriate intervals. Otherwise, gage marks

may be placed on the specimen shoulders, or the overall length

of the specimen may be used to determine elongations (see section

9.6.5); in that case some error is introduced from measurement

across section changes and the results should be qualified.

9.1.2 Reduced Section and Overall Length—Measure the length of the

reduced section (A) and the adjusted length of the reduced section, if

applicable, to the nearest 0.05 mm (0.002 in). If overall specimen length

(L) is used as a basis for determining elongations, measure the initial

overall length of the specimen to the nearest 0.05 mm (0.002 in).

9.1.3 Cross Section—Determine the cross-sectional area of the reduced

section by measurements according to section 7 of Methods E 8 and E 8M.

9.2 Specimen Installation—Leaving sufficient slack for instrumentation

wires so they will not be stretched or crimped during positioning of the

Dewar or subsequent motions during testing, install the specimen in the

cryostat

.

9.3 Seating and Alignment—If the gripping fixture involves loosely

fitting 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 alignment, keep the applied

tensile force below one- third of the proportional limit of the material

being tested. Subsequently maintain a small but sufficient load to ensure

that the alignment is retained during cool-down.

9.4 Cooling Procedure—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. To prevent icing, remove any condensate from the

apparatus before cooling by drying it thoroughly with an air jet or heat

gun. If a clip gage with a protective casing is used, position the gage so

that cryogenic fluid can enter freely to surround the gage's active elements

to prevent the entrapment of gas bubbles and the associated clip gage noise.

314
Next, position the Dewar and precool the apparatus by transferring

liquid nitrogen into the cryostat. After boiling subsides and thermal

equilibrium is reached, 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 8—The heat- transfer characteristics of gaseous

helium are inferior to those of liquid helium; therefore it is

imperative that the specimen remain submerged in liquid helium

to minimize the influence of generated heat on the mechanical

property measurements

.

9.5 Testing Speed:

9.5.1 Rate Control—Since tensile property measurements in liquid

helium are affected by testing speed, the test must include a means of

measuring and controlling the rate of crosshead motion. A nominal strain

rate must be specified, since the actual rate cannot be precisely controlled

or maintained in view of discontinuous yielding. 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.

9.5.2 Rate Limit—Any convenient crosshead speed may be used to reach

an applied stress of one-half the yield strength; after that, the crosshead

speed must be chosen so that the nominal strain rate never exceeds 1 x 10' 3 s' 1

. Higher rates may cause excessive specimen heating and therefore are

not acceptable for basic mechanical property measurements of materials.

9.5.3 Rate Selection—Strain rates ranging from 10" 5 to 10" 3 s* 1 are

generally recommended for tension tests at 4 K, but some alloys are

moderately sensitive to strain rate variations in this range. Some high

strength austenitic steels show mild transitions in tensile properties at

strain rates in the range 10‘ 4 to 10’ 3 s“ l

, and other alloys with high

ratios of strength to thermal conductivity (perhaps titanium alloys) may

show similar trends (15). Consequently, it may be desirable in some tests

to use strain rates much lower than the 1 x 10" 3 s" 1 maximum allowed by this

Method.

9.5.4 Rate Change—It may be desirable to change the strain rate during

a test. For example, the strain required to initiate discontinuous yielding

typically increases with decreasing strain rate. If the first serration

occurs near 0.2% plastic strain, it may be possible to reduce the speed of

the test to postpone the first serration, and to prevent interference in the

measurement of the yield strength (see Fig. 5). This may be accomplished by

first using a relatively low strain rate to determine the yield strength,

and then using a higher strain rate to complete the test.

315
9.6 Measurement of Mechanical Properties:

9.6.1 Force-Extension (Load-Extension) Method—To measure the yield

strength, a record of force 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 by applying the 0.2% offset method, following

Methods E 8 or E 8M. If the 0.2% offset line intersects the curve at a

decrease in force associated with discontinuous yielding, then the highest

stress before that force decrease is reported as the yield strength.

9.6.2 Force-Time Method—Yield strength measurements based on a 0.2%

offset procedure applied to force -versus -time (or force-versus-crosshead)

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 or

some other technique is used, state it clearly in the report.

NOTE 9—Force—time curves for tests at 4 K are typically

nonlinear at the start and less regular than force—extension

curves. The effective modulus of a thermally efficient load

train may be low and dependent on the liquid helium level.

Also the time to achieve a steady state condition may be

affected by changes in heat path efficiency through the

increase of force upon nesting of pull rod interfaces. As a

consequence, yield strength data from force—time curves may be

less accurate than those of the recommended method. For this

reason, check any calculations using time -based deformation

against system stiffness calibrations at the same temperature.

In addition, compare the total system deflection with total

specimen strain using post- test gage mark extension (with

allowance for specimen elastic spring-back, especially if the

material is relatively brittle).

9.6.3 Discontinuous Yielding Stress—Calculate the stress corresponding

to the point of initiation of the first discontinuous -yielding event by

dividing the maximum force sustained at the beginning of the first

measurable serration by the cross-sectional area of the specimen.

9.6.4 Tensile Strength—Calculate the tensile strength by dividing the

maximum force sustained by the specimen during the tension test by the

original cross-sectional area of the specimen.

9.6.5 Elongation—If possible, calculate the percent elongation from

the change in gage length according to Methods E 8 or E 8M. Otherwise,

calculate the percent elongation from the initial (A0 ) and final (Af) values

of the reduced section length: 100- (Af - A0 )/A0 , or from the percent

elongation based on the initial (L0 ) and final (Lf) values of the overall

^ength, and the initial length (A0 ) of reduced section: 100* (Lf - L0 )/A0

.

316
A — Serrations after 0.2% strain

FIG. 5 Stress—Strain Diagram for

by the Offset Method

B - Serrations before 0.2% strain

Determination of Yield Strength

9.6.6 Reduction of Area—Calculate the percent reduction of area

according to Methods E 8 or E 8M.

9.6.7 Rounding Reported Test Data—Round off the calculated numerical

test results according to Methods E 8 or E 8M.

9.6.8 Repiacement Specimens—If necessary, discard any invalid data and

test replacement specimens as per Methods E 8 or E 8M.

10. Report

10.1 Generai:

10.1.1 Materiai Characterization—Describe the test material, including

manufacturing, processing, and metallurgical information.

10.1.2 Specimen Characterization—Describe the specimen location and

its orientation relative to the principal working directions or grain flow

of the stock. Also report the specimen dimensions, including the crosssection dimensions, the fillet radius, the reduced section length, and the

adjusted length of the reduced section (if used).

10.1.3 Strain Rate —

R

eport the crosshead speed and nominal strain rate

for the entire test. If the rate was changed during the test, report the

effective nominal strain rates before and after the rate was changed. If

direct strain measurements were made from the fractured specimen, report the

calculated rates from these data.

317
10.1.4

Mechanical Property Measurements—Report the yield strength, the

tensile strength, the method of offset yield strength measurement, and the

method of extension measurement. Report the region of attachment and the

span of extensometer(s) , if used. Also report the discontinuous yielding

stress and the strain rate at which it was measured, the tensile elongation

and the method of its calculation, the gage- length- to-diameter ratio for

cylindrical specimens, and the reduction of area.

10.2 Optional Data—Report any optional data of importance or interest,

such as measurements of Young's modulus (requires a Class A or Class B1

extensometer) at 4 K, the average grain size of the test material, or the

room temperature mechanical properties, if measured or known. Report the

compliance of the test machine including the cryostat.

10.3 Replicate Tests—If replicate specimens are tested, report the

number of tests, the average values of all mechanical property measurements,

and a measure of the scatter of the data.

10.4 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.

10.5 Anomalies—Report any anomalies in material behavior, test

records, mode of failure, and type and location of fracture.

11. Precision and Bias

11.1 Precision—An overall measure of variability in tensile property

measurements covered in this Method is being established by a series of

interlaboratory tests.

11.2 Bias—There are no accepted standard values for the tensile

properties of materials. In the absence of any true value, no meaningful

statement can be made concerning the bias of data.

12. Keywords

Keywords for this standard are: adiabatic heating, cryogenic properties

(of materials), cryostats, discontinuous yielding, liquid helium, low

temperature tests, mechanical properties (of materials), tensile properties,

tension test.

REFERENCES

(1) 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.

318
( 2 ) Ogata, T. , Ishikava, 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.

(3) Ogata, T. and Ishikawa, K. "Time -dependent Deformation of Austenitic

Stainless Steels at Cryogenic Temperatures," Cryogenics, Vol. 26,

1986, pp. 365-369.

(4) Reed, R. P. , and Walsh, R. P., "Tensile Strain Rate Effects In Liquid

Helium", in: Advances in Cryogenic Engineering

,

Vol. 34, Plenum

Press, 1988, pp. 199-208.

(5) Cryogenics Safety Manual—A Guide to Good Practice, 2nd Ed., Mechanical Engineering Publications, London, 1982.

(6) 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.

(7) 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.

(8) Hartvig, G. and Wuchner, F. , "Low Temperature Mechanical Testing

Machine," Review of Scientific Instruments , Vol. 46, 1975,

pp. 481—485.

(9) 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.

(10) 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.

(11) 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.

(12) Telinde, T. C., "Strain Gages in Cryogenic Environments,"

Experimental Mechanics, Vol. 10, No. 9, 1970, pp. 394-400.

(13) Greenough, R. D., and Underhill, C., "Strain Gauges for the

Measurement of Magnetostriction in the Range 4 K to 300 K," Journal

of Physics E, Vol. 9, 1976, pp. 451-454.

319
( 14 )

(15)

Freynik, H. S., Jr., Roach, D. R. , Deis, D. W. , and Hirzel, D. G.,

"Evaluation of Metal -Foil Strain Gages for Cryogenic Applications in

Magnetic Fields," in: Advances in Cryogenic Engineering , Vol. 24,

Plenum Press, New York, 1978, pp. 473-479.

Reed, R.P., Walsh, R.P., and Tobler, R.L., "Strain-Rate Effects on

Tensile Properties at 4K of a VAMAS Round Robin Austenitic Steel," in

Advances in Cryogenic Engineering , Vol. 36, Plenum Press, New York,

to be published.
Standard-Test Methods-for

Notched Bar Impact Testing of Metallic Materials1

£23: the

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SlmWrnd No. IS lb. Ccmsmi U* DoD Inacx of Sfmzfiauwns ana SianaamJar ittt spmsfle ytm

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mtOud 221.1 ofFtdmwl Tom Mmkod

'of ivutvmtck hat beat u*

1.

Scope

1.1 These test methods describe notched-bar impact

testing of metallic materials by the Charpy (simple-beam)

apparatus and the I»** (cantilever-beam) apparatus. They

give; (a) a description of apparatus. ( b

)

requirements for

inspection and calibration, (c) safety precautions, (d) sampling, (e) dimensions and preparation of specimens. (/)

testing procedures. ( g) precision and bias, and (h) appended

notes on the significance of notched-bar impact testing.

These test methods will in most cases also apply to tests on

unnotched specimens.

if?AThe values stated in SI units are to be regarded as the

standard. Inch-pound units are provided for information

only.

; l&fj'his standard may involve hazardous materials, oper-

;

ations, and equipment. This standard does not purport to i address all ofthe safety problems associated with its use. It is

k

the responsibility of the user of this standard to establish

appropriate safety and health practices and determine the

applicability of regulatory limitations prior to use.

1.2 These methods pertain to tests

at temperatures as low as the boiling

point of liquid nitrogen and higher.

The boiling point of liquid nitrogen

is -196 C (-321 F) at sea level and

is slightly lower at higher elevations. Tests below liquid nitrogen

temperature are more complicated and

i are not covered in this standard.

i

' 1.3 — 1.4

2.

Summary of Test Methods

2.1 The essential features of an impact test are: (a) a

suitable specimen (specimens of several different types are recognized). ( b) an anvil or support on which the test specimen is placed to receive the blow of the moving mass,

ic) a moving mass of known kinetic energy which must be

great enough to break the test specimen placed in its path,

and Id) a device for measuring the energy absorbed by the

broken specimen.

3.

Significance and Use

3. 1 These test methods of impact testing relate specifically to the behavior of metal when subjected to a single application of a load resulting m muldaxial stresses associated with a

notch, coupled with high rates of loading and in some cases with high or low temperatures. For some materials and

temperatures, impact tests on notched specimens have been

found to predict the likelihood of brittle fracture better than

tension tests or other tests used in material specifications.^

Further information on significance appears in Appendix

XI.

For other material-temperature

regimes impact tests of notched

specimens may have less significance .

321
E 23

from its cooiinf (or heating) medium, if used, and positioned on the specimen supports; the pendulum is released without

vibration, and the specimen is broken within 5 s after

removal from the medium. Information is obtained from the

machine and from the broken specimen. The details are described as follows:

1 L2. 1 Temperature of Testing—In most manorial*, impact values vary with temperature. Unless otherwise specified. tests snail be made ax 15 to 32*C (60 to 90*F).^\ccuracy

of results when testing at other temperatures requires the

following procedure: For liquid coding or heating fin a

suitable container, which has a grid raised at least 25 mm (

1

in.) from the bottom, with liquid so that the specimen when

immersed wifi be covered with at least 25 mm (1 in.) of the

liquid. Bring the liquid to the desired temperature by any

convenient method. The device used to measure the temperature of the bath should be placed in the center of a group of

the specimens. Verify all temperature-measuring equipment

at least twice annually. When using a liquid medium, hold

the specimens m an a gitated bath at the desired temperature within ±1*C (±2*F) for at least 5 min. When using a gas

medium, position the specimens so that the gas circulates

around them and hold the gas at the desired temperature

within ±1*C (±2*F) for at least 30 mm. Leave the mechanism used to remove the specimen from the medium m the

medium except when handling the specimens.

Non 8—Temperatures up to +260*C (+500*F) may be obtained

with certain oiis. but “flash-point’' temperatures must be carefully observed.

i* Tests at elevated temperatures or

low temperatures down to liquid nitre

gen temperature may also be performec

according to this standard.

322
appendix

(N«

XI. NOTES ON SIGNIFICANCE OF NOTCHED-BAR IMPACT TESTING

Notch Behavior

11.1 The Charpy V-notch (CVN) impact test has been

msiveiy in mechanical testing of steel products, in

, and in procurement specifications for over three

Where correlations with fracture mechanics paramcjjg available, it is possible to specify CVN toughness

would ensure dasnc-plasnc or plastic behavior

. of fatigue cracked specimens subjected to minimum operating temperatures and maximum in service raxes

of loading.

XI. 12 The notch behavior of the face-centered cubic

metals and alloys, a large group of nonfenous materials and

the austenitic steels can be judged from their common tensile

properties. If they are brittle in tension they will be brittle when notcfaedc whik if they are ducnie in tension they will be

ductile when notched^ aMag^or unusually sharp or deep

notches (much more severe/ than the standard Charpy or

211

Exceptions are known for

some of the face-centered cubic

metals and alloys at extreme

cryogenic temperatures, and

Izod specimens). Even l ew a—pemurea de set akaa thia^

-hnranmmr n f T h mr mairnnn In contrast, the behavior of

the ferrrac steels under notch conditions cannot be predicted

from their properties as revealed by the tension test. For the

study of these materials the Charpy and Izod type tests are accordingly very usefuL Some metals that display normal

ductility in the tension test may nevertheless break in brittle fashion when tested or when used in the notched condition.

Notched conditions include restraints to deformation in

directions perpendicular to the major stress, or multiaxial

stresses, and stress concentrations. It is in this field that the

Charpy and Izod tests prove useful for determining the

susceptibility of a steel to notch-brittle behavior though they

cannot be directly used to appraise the serviceability of a

structure.

323
«N

some ghnawi notch will in most imanrrs place the methods

in their proper order.

XI.4 Teapvatm* Effect

X1.4.1 The taring conditions also affect the notch behavior. So pronounced is the effect of temperamre on the

behavior ofIsted when notched that comparisons are fire- quentiy made by cxaminhn ^penmen fractures and by

. plotting energy value and fracture appearance versos temper-

; axure from tests of notched bats at a series of temperatures When the test temperature has been yarned low enough to

r start deavage hactnre, there may be an extremely sharp drop

f in impact value or there may be a rdanvdy gradual Ming

^off toward the lower temperature. This drop in energy value

;stans when a « begins to exhibit some crystalline

••appearance in the fracture. The transition temperature at £Which this embattling effect takes place varies considerably

'With the size of the pan or test specimen and with the notch

^geometry.

%A

Add a new sac tion XI. 4. 2:

"Section 1.2 restricts the test temperature to no lower f~h»ri the

boiling point of liquid nitrogen. For conventional taats at this

temperature and highar, temperature la simply a parameter which affects

the material performance. At more extreme cryogenic temperatures, the

temperature affects the test dynamics as well as the material

Performance. For example, at liquid hydrogen temperature, —253 °C

(-423°F), or liquid helium temperature, -269°C (-452°F) . tests cannot

ba performed using the procedure outlined in section 11. Metals at

these temperatures have very low specific heats, and any surface

condensation of gases, conductive or convective heat transfer, or Pl**tic deformation quickly produces an extraordinary specimen

temperature rise. Aa a result, (1) specimens cannot bo transferred

from the coolant through air to the anvil in the routine way, and (2) fracture at impact strain rates creates problems of adiabatic heating and temperature control. Various nonstandard methods for Charpy impact tasting below liquid nitrogen temperature ere described in the

H-t«**ture . They incorporate insulating devices and modified

calibration procedures. Initial test temperature as low as -269 °C (-452 °F) may ba achieved using some of those techniques, but

•*treme adiabatic heating on impact la unavoidable and complicates the

d*6* interpretation. At present there ia no standard method for impact

taste belov liquid nitrogen tasqperature , and consensus has not yat bean

reached on the meaning or usefulness of impact data obtained at such

extrasM cryogenic temperatures .

324
U.s. —JAPAN COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

FIRST DRAFT

DECEMBER 1989

NOTE: This is a working document being developed by ASTM subcommittee E

24.08; it is intended for committee review, but not for general

distribution at this time.

A3. TEST PROCEDURES FOR AUSTENITIC ALLOYS IN LIQUID HELIUM AT 4 K

A3 . 1 Scope

A3 . 1 . 1 This annex describes the equipment, procedures, and data analysis used for Jj c tests of austenitic alloys in liquid helium. Ductile

fracture at this temperature involves fast, unstable crack growth. The

procedures set forth in the body of the Standard Method are modified in this

annex to address the specific needs and concerns of cryogenic testing.

NOTE 1 —The boiling point of normal liquid helium (He I)

at sea level is 4.2 K (—452.1 *F or 7.6 8 R). It decreases

slightly with geographic elevation and is 4.0 K (-452. 5 °F or

7.2°R) at the National Institute of Standards and Technology

in Colorado, 1677 m (5500 ft) above sea level. In this

standard the test temperature is designated 4 K.

A3. 1.2 For alloys in liquid helium, Jj c is determined from a pseudoresistance curve plot of J versus Aap. A fatigue -cracked specimen is cooled

in a cryostat to 4 K and tensile force is slowly applied while it is fully

submerged in normal liquid helium (He I). The recommended specimen is a

pin- loaded compact specimen, 25.4 mm (1 in) thick and 50.8 mm (2 in) wide.

The pseudo- resistance-curve slope, AJ/Aap, and the material's apparent

tearing modulus lack their customary significance at this temperature, and

are not measured.

A3 .

2

Definitions

In addition to the information in Section 3, the following definitions

apply:

A3. 2.1 fracture cryostat—an apparatus consisting of a reaction frame

and Dewar for testing specimens in cryogenic environments. A schematic

illustration is shown in Fig. A3 . 1

.

325
LOAD

t

Figure A3.1. Fracture cryostat for tests in liquid helium.

A3. 2.

2

Dewar—a vacuum- insulated container for cryogenic fluids.

A3. 2.

3

adiabatic heating—the internal heating of a specimen caused b testing under conditions such that the heat generated by plastic work

cannot be quickly dissipated from the specimen to the cryogen.

A3 .

3

Summary of Method

A3 . 3 .

1

The object is to estimate the critical value of J near the

initiation of crack growth for austenitic alloys in liquid helium, despite

the occurrence of mechanical and thermal instabilities during the crack

growth stage.

326
A2.3.2 To measure Jjc according to this annex, a fatigue -cracked

specimen is cooled to 4 K in a cryostat containing liquid helium. Using

crosshead, actuator, or clip-gage displacement control, the specimen is

slowly loaded while the load versus load-line displacement curve is

autographically recorded and a plot of J versus Aa is prepared. Data

interpretation and reduction are identical to that specified in the body of

this standard, except that the blunting line approximation for austenitic

alloys at 4 K is J - 4ayAa rather than J - 2oyAa.

A3 . 3 . 3 The compact specimen is preferred for tests in liquid helium

because it offers a large fracture- toughness -measurement capacity for its

size. A relatively low force 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 liquid helium consumption and cost

per test.

A. 3. 3. 4. The single- specimen unloading -compliance procedure (1—3) is

preferred for this environment because, compared to the multiple-specimen

technique, it reduces the required number of specimens, conserves helium,

and lowers the cost of measurements.

A3. 4 Significance and Use

A3. 4.1 Austenitic alloys undergo a change in ductile fracture behavior

at extreme cryogenic temperatures. The mechanism at room temperature is slow

stable tearing, whereas the mechanism at liquid helium temperature consists

of unstable plastic deformation (discontinuous flow) events with incremental

crack extensions. Typically, the load versus displacement record in liquid

helium is composed of continuous and discontinuous deformation stages, as

shown in Fig. A3. 2. At first the deformation is stable with smoothly rising

loads. As testing continues, alternating periods of stable and unstable

loading form a series of serrations on the load versus displacement curves.

At each serration, plastic deformation and tearing occurs abruptly at fast,

uncontrolled rates, and load reduction ensues until the instability is

arrested. The serrations are faint at first but steadily increase in

magnitude. The frequency of serrations and the magnitude of load reductions

is influenced by material and test variables among which are specimen size,

machine stiffness, control mode, and testing rate.

A3.4.2 A constant specimen temperature cannot be maintained at all

times during tests in liquid helium. Due to adiabatic heating, the

temperature in the plastic zone near the crack tip rises above 4 K whenever

the rate of plastic work exceeds the cooling capacity of the cryogen.

Heating during continuous stable loading can be prevented by testing at a slow rate but heating during periods of mechanical instability is inevitable

(4). At slow test speeds, the temperature increase during each instab i lity

is transient and quickly falls back to normal when the instability is

arrested. Typically, altering the test variables can modify but not

eliminate the mechanical instabilities and associated thermal spikes.

327
Q<Zo

Q

O

Figure A3. 2. Typical loading behavior at 4 K, showing serrations.

A3. 4.

3

In principle, the J- integral applies to metals only if they

fracture by slow-stable ductile tearing. Therefore justification for the us of Jj c as a fracture parameter at 4 K rests primarily on the basis of

experimental results. Interlaboratory data for Fe-22Mn-13Cr-5Ni steel show

acceptable agreement of Jjc measurements (5), and intralaboratory data

available for three alloys show little or no variation of Jj c with size for

specimens 12.5 mm (0.5 in) thick or more (6,7). The available data indicatt

that Jj c retains its usefulness as an estimate of fracture toughness near cl onset of ductile fracture in liquid helium, despite discontinuous yielding

events and temporary specimen heating.

A3. 4.

4

By definition, a J-R curve is a plot of resistance to stable

crack extension. The J—versus-Aap plots at 4 K are pseudo -resistance curve

since unstable tearing is involved. The J—versus—Aap plots at 4 K are used

to estimate Jjc by extrapolation, but they lack the physical significance o resistance curves at ambient temperatures where deformation is more nearly

isothermal and crack growth occurs by slow-stable tearing.

A3. 4.

5

The parameter Jjc , operationally defined in this Annex, is

intended primarily for material characterization and may not be directly

applicable to structural designs. Ductile fracture in liquid helium may be

significantly affected by high loading rates, and by the mode of machine

control during testing. Slow loading rates using crosshead, displacement,

clip-gage control are specified in this standard for conventional material

characterization. The effects of severe adiabatic heating at high loading

328
rates must be considered if impact loading is anticipated in service. Also,

the possibility and adverse consequences of unarrested ductile tearing must

be anticipated in the design of structures subjected to load control at 4 K.

A3 . 5 Hazards

Several precautions must be observed when handling cryogenic fluids and

equipment. Skin or eye contact with cryogens will freeze and destroy tissue

(frostbite). The appropriate protection may include goggles, clothing

without pockets or cuffs, gloves, and tongs for handling cold specimens.

Cryogenic containers that are internally pressurized or evacuated are

potentially hazardous in that damage or leaks can produce explosions or

implosions. Also, there is a huge volume expansion when cryogenic liquids

evaporate; therefore asphyxiation is a potential threat if liquid nitrogen or

liquid helium leaks, spills, or is discarded into rooms that are not properly

ventilated. Safety information pertaining to cryogens is available elsewhere

(8). It is the user's responsibility to learn and implement the appropriate

safety practices.

A3 .

6

Apparatus

A3 . 6 .

1

Test Machines—Use a servohydraulic or screw-driven tension test

machine that meets the requirements of Practices E 4 regarding the

verification of force accuracy. The equipment must include a means to

autographically record the applied forces and specimen displacements. Since

the compliance of the test machine influences its response to mechanical

instabilities, it is recommended that the compliance of the test machine and

cryostat at 4 K be measured by replacing the specimen with a rigid block, or

by linking the upper and lower grips without a specimen. The compliance can

then be measured from a plot of force versus crosshead or actuator

displacement up to the highest force used to qualify the machine.

A3. 6.

2

Cryogenic System Design—Since most conventional test machines

have a maximum force of 10 tons (22 480 lbf)

,

it is recommended that the

apparatus be designed to accommodate a standard, 25.4 mm (1 in) thick compact

specimen (see Section 8.1.3). A servohydraulic test machine with an inverted

stage (actuator mounted above the test chamber) and a reaction frame that is

free at the lower end is recommended. These features are desirable for

operating convenience and cost effectiveness for testing at 4 K (9)

.

A3. 6.

3

Cryogenic Construction Materials—Many structural materials,

including the vast majority of ferritic steels, are brittle at 4 K.

Therefore cryostats must be well designed and fabricated using cryogenic

materials of adequate strength and toughness. Materials that have low

thermal conductivity are favored to reduce heat conduction. Austenitic

stainless steels are widely used, especially those strengthened by the

addition of 0.05 to 0.4 weight percent nitrogen. Maraging steels plated with

nickel to prevent rusting have been used for specimen clevises. High purity

alpha—titanium alloys are especially appropriate for some applications

requiring high strength- to -thermal conductivity ratios, such as pull rods.

329
Nonmetallic materials (for example, fiber -reinforced composites) are

excellent insulators and have been used for compression members in some

cryostats

.

A3. 6 . 4 Cryostats and Support Apparatus:

A3. 6. 4.1 Cryostats—A cryostat capable of retaining liquid helium for

the duration of testing is required. Cryostats are usually custom-built, but

they may be designed to accommodate commercially available Dewars. Efficient

designs consider thermal and mechanical trade-offs (9). Several designs are

possible, including turret-disc devices for sequential testing of several

specimens with a single cooling (3,7,9—11).

A3. 6 . 4. 2 Dewars—Stainless steel Dewars are recommended for low maintenance, durability, and safety. Glass Dewars are fragile and hazardous if

broken, whereas fiberglass Dewars are relatively expensive and may require

maintenance. A single Dewar is sufficient for tests of short duration. The

alternative is a double -Dewar arrangement using an outer Dewar containing

liquid nitrogen to surround the helium Dewar for greater thermal efficiency.

A3. 6 . 4. 3 Ancillary Equipment—Dewars and transfer lines for handling

liquid helium must be vacuum insulated. Pressurized gases, vacuum pumps, and

liquid nitrogen apparatus are therefore necessary. After testing, the helium

may be recycled as a gas or liquid, or it may be discarded to the atmosphere.

Recycling as gas or reliquefaction requires a substantial investment in

purification and cryogenic systems.

A3. 6 . 5 Grips and Pins—Conventional grips conforming to the design in

Fig. A2.2 may be used, but alloys with adequate 4-K mechanical properties are required for construction. Pins and grips require adequate 4-K strength,

fracture toughness, and fatigue resistance. Pins should be manufactured from

high-strength materials to resist bending. Maraging steels can be used for

grips of the conventional design. Lower strength materials such as AISI

304LN may be satisfactory for grips if the clevis arm thickness is enlarged.

A3 . 6 . 6 Extensometer :

A3. 6. 6.1 Requirements—In addition to rendering accurate displacement

measurements at 4 K, cryogenic extensometers must sustain thermal shock and

moisture condensation associated with thermal cycling between room and liquid

helium temperature. Conventional clip gages that operate in liquid helium

may be purchased or built using cryogenic materials and fabrication

techniques. Ring-shaped clip gages were originally designed for cryogenic

tests, and they offer some practical advantages for tests in liquid helium

( 12 ).

A3. 6 . 6 . 2 Construction—Select a satisfactory combination of active

strain gage elements, backing materials, and bonding agents for cryogenic

extensometers. One common choice is a Ni—Cr-alloy gage with a temperaturecompensated active element. Typically, four 350-0 gages are wired in a full

Wheatstone bridge. Leadwires covered with polytetrafluoroethylene insulation

and soft solders having low tin contents are recommended to resist mechanical
failures and electrical shorts induced by thermal cycling between room

temperature and 4 K.

A3.6.6.3 Calibration—

C

alibrate the clip gage or other displacement

measuring device at room temperature and in liquid helium. This may be done

with a barrel-micrometer fitted with vertical extension tubes, with the gage

mounted at the lower end and immersed in liquid helium for calibration at 4 K

(12). Once the calibration is known and proven linear, accurate, and

reproducible, then room- temperature checks will suffice to verify the 4-K

calibration indirectly. Nevertheless, the calibration should be repeated

periodically at 4 K, especially when damage is suspected or repair is

performed.

A3.6.6.4 Temperature Maintenance and Liquid-Level Indicators—The

intended cryogenic bath temperature of 4 K is maintained by the fixed boiling

point of liquid helium. It is therefore unnecessary to monitor the specimen

temperature directly. Instead, a suitable meter or indicator is required to

ensure the proper liquid level. Use an on—off carbon- resistor meter, a

continuous -level meter with superconducting wires, or a thermocouple (Chromel

versus gold + 0.07% iron which offers maximum sensitivity at temperatures

near absolute zero) . Position the indicator in the cryostat above the

specimen and monitor the liquid level to ensure that the specimen remains

submerged throughout the test.

A3. 7 Sampling

Remove samples for specimen preparation from the material in its final

metallurgical condition. Perform any heat treatments before sampling and

specimen machining. Choose the sample location to be representative of the

stock material, realizing that samples from one portion of the stock may not

be representative of the whole, and allow for areas affected by the cutting

process. The usual two- letter specimen orientation code of Methods E 399-83

Section 5.1.3 applies.

A3.

8

Specimens

A3. 8.1 Specimen Configuration:

A3. 8. 1.1 Specimen Types—Various specimen configurations might be used,

but the compact specimen offers strong advantages and is preferable to other

specimens for routine tests in liquid helium. Bend specimens, for example,

require greater mass for an equivalent ligament size, and the long span of

the bend specimen transverse to the load axis requires fixtures that are

inconvenient for some cryostats.

A3 . 8 . 1 .

2

Standard 4 -K Test Specimen—A standard compact specimen

geometry and size is specified for material evaluations at 4 K. The standard

specimen has a thickness, B, of 25.4 mm (1.0 in) and a width, W, of 50.8 mm

(2.0 in). The W/B ratio is 2.0, the planar dimensions are proportional to W,

and the starter notch is according to the guidelines of Section 7.3. This

331
specimen is large enough to satisfy the size requirements for most austenitic

alloys tested at 4 K.

A3. 8. 1.3 Alternative Specimens—If the standard specimen geometry or

size is not practical for any reason, select an alternative. Factors

toconsider are: the technical purpose, the load limit of the test machine,

the stock dimensions, and the need for representative samples of the test

material

.

A3. 8.

2

Fatigue-Precracking:

Fatigue-precracking should normally be conducted with the specimen in

liquid helium or liquid nitrogen. However, some test materials may be

precracked in air at room temperature if it can be demonstrated that room

temperature precracking does not affect the toughness subsequently measured

at 4 K. For the precracking to be valid, the criteria of Section 7.6 must

be satisfied for the temperature selected. If precracking is performed at

liquid helium temperature, use 4-K values of oy in 7.6.1 and E in 7.6.2. If

precracking is performed in liquid nitrogen or at room temperature, use the

appropriate cry and E values for these temperatures

.

NOTE 2—The purpose of precracking at liquid nitrogen

or room temperature is to conserve liquid helium. The

boiling point of nitrogen at sea level is 77 K (-321°F, or

138. 7°R)

.

For most alloys, precracking at this

temperature is virtually equivalent to precracking in

liquid helium, and it is convenient since 4-K test

specimens are routinely precooled in liquid nitrogen

anyway. On the other hand, precracking at room

temperature toughens some (not all) alloys at 4 K.

Therefore the precracking of 4-K test specimens at room

temperature is acceptable only if it can be demonstrated

for the specific alloy being tested that the Jjc value at

4 K is not affected.

A3.8.3 Side Grooves—Side grooves might hinder or enhance valid

measurements at 4 K, and their use is an option to be considered case by

case. At room temperature, side grooves are sometimes employed to alleviate

fatigue crack- front curvature and crack- front tunneling, or to improve the

agreement between the optically measured and compliance -calculated end-oftest Aa values. Side grooves may serve the same purpose at liquid helium

temperature; however, side grooves can lower the Jj c measurements for some

materials, and they always promote larger serrations at 4 K. The latter

effect can be conducive to invalid results for some materials by making it

more difficult to satisfy the Aap distribution requirements.

A3 . 9 General Procedures

A3. 9.1 Specimen Installation—Install the specimen in the cryostat,

leaving sufficient slack for the instrumentation wires so that they will be

332
neither stretched nor crimped during positioning of the Dewar and subsequent

testing. Before cooling, check the seating of the specimen and grips at room

temperature. The specimen may be preloaded slightly for this purpose.

A3 . 9 .

2

Cooling—Ice that forms within the cryostat during cool-down can

block cryogenic transfer lines or cause erratic loading behavior. Prevent

icing by thoroughly drying the clip gage, specimen, grips, loading pins, and

load train members thoroughly before cool-down with a pressurized air jet or

heating gun. Also, purge any condensate from the transfer lines. Position

the Dewar and precool the apparatus by transferring liquid nitrogen into the

cryostat. After the specimen reaches thermal equilibrium, remove all of the

liquid nitrogen and add liquid helium to the cryostat. Add extra liquid

helium to compensate for evaporation that will occur 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 3—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 from plastic work during the test.

A3. 9. 3 Loading:

A3. 9. 3.1 Control Mode—Load the specimen using crosshead, actuator, or

clip- gage control. The mode of control determines the machine's response

during instabilities and influences the frequency and shape of serrations

formed on the force—displacement records.

NOTE 4—For screw- driven machines where crosshead speed

is controlled, and for servohydraulic machines where

actuator position is controlled, the clip gage opens as

the load drops with each serration. For servohydraulic

machines in clip gage displacement control, the gage

displacement is maintained constant by the machine as the

load drops with each serration. Clip gage displacement

control arrests the instability events sooner than

crosshead or actuator control. Thus clip gage control

produces smaller, more numerous serrations which benefits

some tests by affording more opportunities to locate J-Aap

data in the exclusion interval.

A3. 9. 3.

2

Testing Rate—Tests in liquid helium must be performed slowly

to minimize specimen heating. Load the specimen at a rate such that the time

to reach the load 0.4Pl is between 1 and 10 min. Higher rates may cause

excessive heating and obscure the material property characterization.

NOTE 5—The testing rate should be slow enough to

prevent a temperature rise of more than 1 kelvin during

the initial period of smoothly rising loads; it should

333
also be slow enough to allow the specimen temperature to

return to 4 K after the arrest of each instability.

Heating is unlikely to present a practical problem in

single specimen tests that are periodically interrupted

for unloading compliance measurements . Of greater

potential concern are multiple specimen tests where

specimens are loaded without interruption to significant

plastic displacements. The heating that occurs in those

tests is a function of material and test variables; it

increases with testing rate, specimen ligament size, and

the plastic component of J associated with fracture, Jp,

and is favored by the low specific heats and thermal

properties of materials at extreme cryogenic temperatures

.

Few quantitative measurements of specimen temperature rise

have been made, but observations for AISI 310 S steel (Jjc - 333 kJ/m2

, or 1900 in’ lb ’in" 2

) may be taken for a

guideline; for a standard specimen (5 - 25.4 mm, W - 50.8

mm, and a/W - 0.7) tested in liquid helium, the internal

heating during the initial continuous loading and the

periods of stable loading between serrations is virtually

eliminated using crosshead speeds of 0.2 mm/min (0.008

in/min) or less

.

A3. 9. 4 Data Collection—During periods of instability it is impossible

to predict or control the extent of incremental crack growth. Opportunities

for unloading and Aap determination exist only during the intervals of stable

loading. Thus, the number and size of serrations influences the J—Aap data

distribution that can be obtained in the exclusion interval, and there is no

guarrantee that the data distribution requirements of Section 9.2.2 can be

satisfied for a given alloy and specimen size.

NOTE 6—Large serrations are problemmatic for some

material—test conditions. In some cases, if it is

difficult to obtain the required number or spacing of

J—Aap points, it may be possible to meet the requirements

of Section 9.2.2 by: (1) lowering the machine compliance

by stiffening the load train, (2) switching to a smaller

or thinner specimen, (3) switching from side -grooved to

ungrooved specimens, or (4) switching from crosshead to

clip- gage displacement control.

A3. 9. 4.1 Single Specimen Tests—Unloadings for Aap determination may

be chosen randomly or at regular intervals at the experimenter's discretion.

Unloadings performed just before load point discontinuities tend to produce

J-Aap pairs at the upper bound of pseudo -resistance curve scatter, whereas

those at arrest points just after load point discontinuities tend to produce

J—Aa

p pairs at the lower bound of scatter. This effect becomes significant

if tne serrations are very pronounced. In that case, the following approach

might be used to reduce the data scatter: during the initial stage of

testing, while the load is smoothly rising, unloadings might be performed

anytime; once serrations begin, the apparent scatter of J-Aap might be

334
reduced by performing the unloadings preferentially at the natural arrest

points after serrations.

A3. 9. 4. 2 Multiple Specimen Tests—Similarly, in multiple specimen

tests, the scatter of J-Aap plots might be reduced if each specimen loaded

into the regime where serrations are formed is unloaded at a natural arrest

point just after a serration.

A3. 10 Interpretation and Calculation of Results

A3. 10.1 Calculate the area under the load-versus- displacement curve,

including any area increase associated with serrations.

NOTE 7—For tests in clip-gage displacement control,

the area does not increase during the period of

discontinuous flow. Tests in crosshead displacement

control give higher values of J for a given crack

extension, owing to the fact that the area does increase

with serrations. In either case, the data are acceptable:

the intent is not to determine a physically meaningful

J—Aap curve, but merely to obtain a systematic data

distribution permitting extrapolation to Jj c

.

A3. 10. 2 Construct the blunting line (assuming J - 4<7yAa) . Then plot the

regression line, and qualify the J—Aap calculations following Section 9.2.

A3. 11 Report

A3. 11.1 Material Characterization—Describe the material tested and

report its yield strength, tensile strength, and Young's elastic modulus at

room temperature and at 4 K. Supplementary information would include a

desription of the material product form, processing, heat treatment, and any

other relevant metallurgical information.

A3. 11. 2 Fracture Toughness Measurements—Report each Ji c determination

in units of kJ/m2 (in* lb* in" 2

) and assess the validity according to Section

9.2. State the number of replicate tests performed, the individual and

average Jjc values, and a measure of scatter.

A3. 11. 3 Specimen Characterization—Report the crack-plane orientation

and the specimen dimensions, including thickness B, width W, and initial

ligament size b 0

. For side-grooved specimens, report Bfj and any other

supplemental information about the side -groove geometry.

A3. 11. 4 Fatigue-Precracking Conditions—Report the temperature of

fatigue -precracking, the maximum fatigue load, and the maximum stress

intensity factor or the stress intensity factor range. Supplemental

information would include the final fatigue crack length obtained using the

nine point average, the measured surface crack lengths, and the maximum

deviation of any single measurement from the average.

335
A3. 11. 5 Test Conditions—Report the testing rate in mm/min (in/mm) and

specify whether crosshead, actuator, or clip -gage opening displacement was

controlled. Supplementary information may include the compliance of the test

apparatus

.

A3. 11.6 Fracture Surface Appearance and Anomalies—Describe the

fracture -surface appearance and report any anomalies in material behavior,

test records, or failure mechanisms.

A3 . 12 Accuracy and Precision

A3. 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

.

A3. 12.2 Precision—Information on the precision of 4-K fracture tests is

being obtained from interlaboratory test programs currently in progress.

A3. 13 Keywords

Keywords for this annex are: adiabatic heating, cryogenic properties (of

materials), cryogenic tests, fracture toughness, liquid helium, and low

temperature tests.

REFERENCES

(1) 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.

(2) 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.

(3) Shimada, M. , Ogawa, R. , Moriyama, T. , and Horiuchi, T. , "Development

of a Cryogenic Fracture Toughness Test System," Cryogenic Technology

,

Vol. 21, No. 5, 1986, pp. 269-274.

(4) Tobler, R. L. , Shoji, T., Takahashi, H. , and Ohnishi, K. , "Fracture,

Acoustic Emission, and Adiabatic Heating of Austenitic Stainless

Steels at Liquid Helium Temperature," in: Progress in Acoustic

Emission III, The Japanese Society of NDI, Tokyo, 1986, pp. 453-461.

336
(5) 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—22Mrt-13Cr-5Ni Austenitic Stainless Steel at 4 K," in:

Advances in Cryogenic Engineering—Materials

,

Vol. 34, Plenum Press,

New York, 1988, pp. 241—249.

(6) Shimada, M. , Tobler, R. L. , Shoji, T. , and Takahashi, H. , "Size, SideGrooving, and Precracking Effects on Jj c 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.

(7) Yoshida, K. , Nakajima, H. , Oshikiri, M. , Tobler, R. L. , Shimamoto, S.,

Miura, R. , and Ishizaka, J., "Mechanical Tests of Large Specimens at

4 K: Facilities and Results," in: Advances in Cryogenic Engineering—

Materials

,

Vol. 34, Plenum Press, New York, 1988, pp. 225-232.

(8) Cryogenics Safety Manual—A Guide to Good Practice, 2nd Ed. , Mechanical Engineering Publications, London, 1982.

(9) 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.

(10) 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.

(11) 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.

(12) 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.

337
-
TECHNOLOGY TRANSFER

339

TECHNOLOGY TRANSFER PROGRAM

LEADER: N.J. Simon

CONTRIBUTORS: E.S. Drexier, R.P. Reed

OBJECTIVES

o Organization of and participation in 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 low- temperature data for structural alloys, composites, and

weldments

.

o Preparation of an annual report to the Office of Fusion Energy, U.S.

Department of Energy.

o Evaluation of low- temperature mechanical and physical properties of

insulator, conductor, and structural materials for cryogenic copper and

superconducting magnets; preparation of handbook pages and supporting

documentation; distribution of handbook pages to participants in the

fusion energy program and to the Materials Handbook for Fusion Energy

Systems

.

RESEARCH HIGHLIGHTS:

o The Twelfth Annual Cryogenic Structural Materials Workshop was

organized; it was held October 11-12, 1989 in Vail, Colorado.

o "Materials Studies for Magnetic Fusion Energy Applications - XII,"

(NISTIR 89-3931, 277 pages, 1989, editors R.P. Reed and R.L. Tobler)

was prepared, published, and distributed.

o Handbook pages on the cryogenic tensile, elastic, electromagnetic,

fatigue, and thermal properties of C17000-C17510 beryllium- copper were

prepared and distributed.

o Handbook pages on C50500—C52400 phosphor-bronze tensile, elastic,

fatigue, electromagnetic, and thermal properties are in preparation.

Documents (over 250) have been acquired, and preliminary data sets

extracted and analyzed. Equations relating yield strength, tensile

strength, and elongation to cold work, temperature, Sn content, and

grain size between 4 K and room temperature have been developed from

the data base. Further analysis of other properties is in progress.

These alloys are used as stabilizers for Nb 3 Sn superconducting cable.

o The cryogenic data base for AISI 316LN has been updated with data

acquired since the AISI 316 handbook pages were prepared in 1982. A

preliminary report on this data base was presented at the ITER

Specialists Meeting on Magnet Materials in Garching in February, 1990.

Alloy 316LN is a prime candidate for use in thick sections in the

toroidal field coils in the ITER design.

341
'
TWELFTH CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

Vail, Colorado

October 11-12, 1989

The 12th in this series of workshops was attended by about 40 participants.

NIST-Boulder staff made the program arrangements, presented nine invited

talks, and led three discussion sessions. Other participants were from DOE

fusion energy and accelerator projects, national laboratories, universities,

and industry. The workshop provided an opportunity for discussion of

research accomplishments, problems, and goals in the development of

materials for cryogenic and superconducting magnet systems for fusion

reactors and high-energy accelerators. The following topics were

highlights

o Materials issues for ITER (International Thermonuclear

Experimental Reactor) and CIT (Compact Ignition Tokamak) including

short crack fatigue and fracture toughness

.

o Manufacture and design of radiation-resistant 3-D composite

insulators

.

o Development of 4-K mechanical test standards for structural alloys

and composites.

o International Collaborative Programs (with USSR, Japan, and

Europe

)

o Development of new structural design codes for cryogenic

temperatures

.

A copy of the program and a summary of the technical presentations follow.

In a few cases, there were changes in the program as noted in the summary of

the workshop.

343
12th CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

PROGRAM

WEDNESDAY. OCTOBER 11

INTRODUCTIONS

8 : 30A Welcome

Office of Fusion Energy Overview

NORMAL-METAL CONDUCTOR MAGNETS

CIT Magnet Design - Structural Alloy Requirements

CIT Materials, Issues, and Results from

Conductor Research Program

Type 718 Draw Bar Structure Problems for C-MOD

Creep of Copper: 4 to 295 K

COLLABORATIVE PROGRAMS

Discussion: US/Japan and (possible) US/USSR Joint

Programs on Low Temperature Structural Materials

LUNCH BREAK

STRUCTURAL INSULATORS

1:30P CIT Insulation Issues and Research Program Results

Proton Irradiation Study of Insulating Materials:

Preliminary Results

Applications of Composite Materials in Superconducting

Particle Accelerator Magnets

Manufacture of New 3D Composite Insulators

Development of 3D Composites

Current Status of Interlaminar Shear Testing of

Composite Materials at Cryogenic Temperatures

DESIGN CODES

Discussion: Structural Design Codes

for Superconducting Magnets (Leader)

R. Reed, NIST

D. Beard, OFE

J. Citrolo, PPPL

C. Bushnell, PPPL

H. Becker,

R. Ballinger,

I . Hwang , MIT

N. Simon, NIST

R. Reed, NIST

(Leader)

T . McManamy , ORNL

N. Munshi, CTD

T. Nicol, Fermilab

N. Munshi, CTD

J. Benzinger,

Spaulding

M. Kasen, CTD

H. Becker, MIT

5 : 30P RECEPTION (Altitude Club)

344
THURSDAY. OCTOBER 12

SC MAGNET STRUCTURAL ALLOYS

8:30A ITER Material Issues L. Summers, LLLN

Fatigue Problems of ITER Central R. Reed, NIST

Solenoid Coils and Other Topics

Properties of Austenitic Stainless Steels J. Chan, U. Cal

in High Magnetic Fields

Progress Report: Handbook of Structural Materials N. Simon, NIST

for Superconducting Magnets

Fatigue of Metastable Austenitic Stainless Steels M. Xequn and

J. Morris, U. of Cal

Cryogenic Properties of Incoloy 908

Research on Cryogenic Properties of Al-Li Alloys

Large Scale Cryogenic Test Facility at NIST

LUNCH BREAK

STRUCTURAL ALLOY WELDING

1:30P Mechanical Properties of Low-Ferrite Welds at 4 K

Weld Alloy Design

Welding Problems, Panel Discussion

I. Hwang, M. Morra,

R. Ballinger, MIT

J. Morris, U.Cal

D. McColskey , NIST

C. McCowan, NIST

T. Siewert, NIST

T. Siewert, NIST

(Leader)

METAL AND COMPOSITE TEST STANDARDS

Discussion: Low Temperature Mechanical R. Tobler, NIST

Property Test Standards: Structural Alloys and Composites (Leader)

The Effect of Strain Rate on Low-Temperature, R. Walsh, NIST

Round-Robin Tensile Tests

Warm Prestress Effects in Cryogenic Structural Steels R. Tobler, NIST

J- Integral Test Analysis Procedures and Interpretation P. Purtscher, NIST

345
SUMMARY OF TECHNICAL PRESENTATIONS

INTRODUCTION

Office of Fusion Energy Overview : D. Beard, Office of Fusion Energy,

Department of Energy, Germantown, MD

The importance of cryogenic properties of structural materials and

electrical insulators to the design of high field superconducting magnets

design was emphasized. The next generation of high field magnets for plasma

confinement may require Nb 3 Sn rather than the currently used NbTi as

superconductors. Model or prototype coil test facilities are necessary to

develop magnets with this new conductor. Conductor test facilities such as

FENIX at Lawrence Livermore National Laboratory (LLNL) are also needed.

Fusion energy systems being developed in Japan, Europe, and the Soviet Union

were reviewed. International and U.S. task- sharing must be worked out so

that progress can be made with available funding.

NORMAL-METAL CONDUCTOR MAGNETS

CIT Magnet Design -Structured Alloy Requirements : J. Citrolo, Princeton

Plasma Physics Laboratory, Princeton, NJ

An all -welded structure has been chosen for the toroidal field (TF)

coils of the Ignition Tokamak (CIT) . The choice of welding was made because

the inner legs of the TF coil cases must be joined to a complex bucking

cylinder to resist strong overturning moments. The presentation generated

discussion about the use of 21-6-9 as the baseline alloy for the structure,

because of poor reports of weld properties and low cryogenic fracture

toughness. This alloy does have very high strength compared to 316LN.

Further studies on weldability are planned.

CIT Materials. Issues, and Results From Conductor Research Program : C. Bushnell, Princeton Plasma Physics Laboratory, Princeton, NJ

With the change from a wedged to a bucked design for the TF coil

systems, conductor requirements have also changed somewhat. Currently,

three systems are undergoing investigation.

(1)

A copper- Inconel alloy laminate with less Inconel than in previous

designs

(2)

Copper Cl 75 10 that has been ove raged to provide a high

conductivity combined with yield strengths in the

range of 550-620 MPa.

(3)

Copper C15715, with lower yield strength but higher

conductivity.

346
Alloy 718 Drawbar Structure Problems for Alcator C-Mod : R. Balllinger, H.

Becker, I. Hwang, and M. Morro, Massachusetts Institute of Technology,

Cambridge , MA

In the TF magnet of the C-Mod fusion energy system, the critical

parameter for the drawbars that withstand the 110 MN separating force is

fatigue life. Crack growth and fracture toughness control the permissible

initial crack size. Low fracture toughness was observed after aging of the

forged and machined bar to increase yield strength. Therefore, the bars

were given an annealing treatment, which lowered the yield strength, but

improved the predicted life.

Creep of Copper: 4 to 295 K : N.J. Simon, R.P. Walsh, E.S. Drexler, and

R.P. Reed, National Institute of Standards and Technology, Boulder, CO

Creep of copper is of interest both for the CIT machine and for

superconductoring magnets, where copper is used as a stabilizer.

Measurements of creep were made at 4, 76, and 295 K for periods of several

months. Results were compared with logarithmic and steady state creep

models

.

COLLABORATIVE PROGRAMS

Discussion: U.S. /Japan and (possible) U.S./USSR Joint Programs on Low

Temperature Structural Materials : R.P. Reed (Leader), National Institute of

Standards and Technology, Boulder, CO

Cooperative programs to support the design issues of ITER (International Thermonuclear Experimental Reactor) were the chief topics of

discussion. Development of a short -crack fatigue database on the candidate

alloys for the conductor jacket of TF coils was one issue identified.

Another issue was the writing of specifications for ITER materials such as

the structural alloys to be used in very thick (-100-mm) sections in the TF

magnet case. Development and adequate cryogenic testing of radiationresistant insulators was another topic of discussion. A Japanese company

has developed a 3-dimensional (3-D) fiberglass weave that is under

consideration for use in the magnet insulation. This or another material

would be wound around the conductor assembly and vacuum- impregnated with a

resin. Creep of such a composite insulator as well as radiation damage to

it are possible failure modes for the magnets.

STRUCTURAL INSULATORS

CIT Insulation Issues and Research Program Results : T . J

.

McManamy , Oak

Ridge National Laboratory, Oak Ridge, TN

The change in the CIT design described above also modified the

requirements for the insulation. Stresses were considerably reduced, but an

additional requirement of a shear strength of 21 MPa with no compression was

347
generated. In the research program, irradiation and testing at room

temperature were carried out on two bismaleimide systems and one epoxy resin

system. Two of the systems used a Japanese 3-D weave. Cryogenic thermal

cycling of insulation systems was also investigated.

Proton Irradiation Study of Insulating Materials: Preliminary Results . N.A. Munshi, CTD, Inc., Boulder, CO

A linear accelerator at Brookhaven National Laboratory was used to

study the effect of proton irradiation on nine prototype magnet insulating

materials. Considerable thermal damage was sustained by the specimens,

which made it difficult to draw clear conclusions. However, one of the

epoxy resin systems with a long pot life showed promise in the tests. A

resin system with a long pot life is very desirable for the lengthy process

of vacuum impregnation of a large superconducting magnet.

Applications of Composite Materials in Superconducting Particle

Accelerator Magnets : T.H. Nicol, Fermi National Accelerator Laboratory,

Batavia, IL

Use of composite materials for electrical insulation and as thermal

standoffs in structural applications was reviewed. The Tevatron accelerator

at the Fermi National Accelerator Laboratory has provided experience with

the operation of many superconducting magnets over an extended time period.

A sizeable fraction of the magnets have had to be repaired or replaced,

sometimes due to inadequate composite properties under extended periods of

use. The low shear modulus of structural composites is also a potential

problem when transporting a large structure such as a superconducting

magnet.

Manufacture of New 3-D Composite Insulators : N.A. Munshi , CTD , Inc .

,

Boulder, CO

Procedures for vacuum- impregnation fabrication of specimens were

presented. The methods developed permit production of void- free laminate

specimens reinforced with glass. Tests on specimens produced with these

methods will be used to model conditions encountered with superconductor

magnet insulation.

Development of 3-D Composites : J.R. Benzinger, Spaulding Composites

Company, Inc., Tonawanda, NY

The extent of commercial development and availability of 3-D fabrics

for composites was reviewed. Recent developments of these materials in

Japan were described. No off-the-shelf 3-D materials are currently

available in the U.S. It is hoped that the use of such designs will lead to

improvements in interlaminar shear strength of composites.

348
Current Status, of Interlaminar Shear Testing of Composite Materials

at Cryogenic Temperatures : M.B. Kasen, CTD, Inc., Boulder, CO

Methods of interlaminar shear measurement were illustrated. All such

test methods have the objective of producing pure shear over a plane.

However, all methods introduce some tensile or compressive components in the

stress state. Tensile stresses reduce the measured shear strength below its

true value, while compressive stresses increase it. A further difficulty is

the requirement to miniaturize the tests for cryogenic use. A variant of a

V-notch shear test may prove adaptable for tests at cryogenic temperatures.

Discussion: Materials Questions for the Superconducting Super Collider

(SSC) and Continuous Electron Beam Accelerator Facility (CEBAF) : J. O'Meara, CEBAF, Newport News, VA; W.A. Smith, SSC Laboratory, Dallas, TX;

T.H. Nicol, Fermi National Accelerator Laboratory, Batavia, IL

An informal discussion summarized materials research needs of

accelerator projects using superconducting magnets. Needs identified were:

(1) radiation resistance of composites, (2) creep of conductors and

composites, (3) fatigue of composite support materials during transport, and

of thin-wall stainless steel bellows, (4) methods of weld inspection, (5)

thermal conductivity data on structural composites, (6) development of

"standard" composites.

DESIGN CODES

Discussion: Structural Design Codes for Superconducting Magnets : H. Becker, (Leader), Massachusetts Institute of Technology, Cambridge, MA

It was pointed out that previously developed structural design codes

relied on failure data, but such data do not exist for superconducting

magnets in fusion reactors. Consequently, codes such as the ASME Boiler

Pressure Vessel Code (BPVC) have been used as starting points. The BPVC

specifies design allowable stresses such as the lesser of 1/3 of ultimate

strength or 2/3 of yield strength. This code was also based on periodic

inspection of the pressure vessel structure, but inspection during operation

will usually be impossible for the magnets in a fusion reactor.

SUPERCONDUCTING MAGNET STRUCTURAL ALLOYS

ITER Materials Issues : L.T. Summers, Lawrence Livermore National

Laboratory, Livermore, CA

The design of ITER magnets will place challenging requirements on

materials. A reaction heat treatment at temperatures up to about 725°C will

be applied after the forced flow conductor system is fabricated.

Consequently, alloys for the conductor sheath have to be evaluated after

this type of heat treatment. A major issue is the fatigue crack growth rate

of short cracks (-150 pm) in these alloys. The external magnet cases

require thick section plate that poses development and procurement problems

as well as potential problems in welding. The radiation dose that the

magnet insulation must sustain is estimated to be 1—2 x 10 Gy.

349
Fatigue Problems of ITER Central Solenoid Coils and Other Topics : R.P.

Reed, National Institute of Standards and Technology, Boulder, CO

There is a lack of measurements of fatigue crack growth rate in the

short crack (-100 pm) region. A new measurement technique using a variable

ratio of minimum to maximum stress intensity factor can be used to supply

data in this domain. The number of cycles of fatigue to which the central

solenoid will be subjected, - 200,000, is close to the expected lifetime of

proposed conductor sheath alloys such as Incoloy 908 and JBK-75.

Fracture of Metastable Austenitic Stainless Steels in a High Magnetic Field

Environment : J.W. Chan and J.W. Morris, Jr., Center for Advanced Materials,

Lawrence Berkeley Laboratory, Berkeley, CA

Metastable austenitic stainless steels are potential structural alloys

for high-field superconducting magnets. Previous work showed an increase in

fracture toughness of AISI 304 in an 8-T magnetic field. The tests on AISI

304 were repeated and additional tests were run on AISI 310S, stable with

respect to transformation at 4 K, and AISI 304L, less stable than the AISI

304 tested. An increase in fracture toughness in an 8-T field was found for

the stable AISI 310S, and a decrease was observed for the relatively

unstable 304L. These results were contrary to what was expected from

previous work and will be investigated further.

Progress Report: Handbook of Structural Materials for Superconducting

Magnets: N.J. Simon, E.S. Drexler, R.P. Reed, National Institute of

Standards and Technology, Boulder, CO

This handbook differs from earlier handbooks in that much more

documentation of design curves and test measurements is presented. The

materials covered are those to be used in fusion applications and priorities

are set by guidance from fusion materials workshops. To date, materials

covered have included AISI 316 and 316LN, 21-6-9, AISI 304LN, 304HN, and

304LHN, C10100-C10700 coppers, and C17000-C17510 beryllium coppers. Work on

Cu-Sn alloys, stabilizers for Nb3 Sn superconductors, is currently in

progress. Future plans are to update the AISI 316LN database with recently

acquired NIST and European data, and to acquire data on the alloys proposed

for the cable- in-conduit conductor (CICC) configuration.

Fatigue Crack Growth in Metastable Austenitic Stainless Steels : M. Xequn

and J.W. Morris, Jr., Center for Advanced Materials, Lawrence Berkeley

Laboratory, Berkeley, CA

AISI 304L, a metastable austenitic stainless steel, has a lower fatigue

crack growth rate at 77 K than the more stable AISI 304LN. The hypothesis

was advanced that transformation during fatigue induces crack closure,

although the volume expansion of the transformation could tend to prop the

crack open. Near the threshold of fatigue crack growth, closure dominates.

Transformation effects on the fatigue crack growth rate are very

complicated, because 21-6-9, with considerable strain- induced martensite,

has a relatively high fatigue crack growth rate.

350
Ballinger, M.M. Steeves, and M.O. Hoenig, Massachusetts Institute of

Technology, Department of Materials Science and Engineering, Plasma Fusion

Center, Cambridge, MA.

The alloy Incoloy 908 was developed specifically for use with Nb 3 Sn

superconductors in the CICC configuration. Weldability, maintenance of good

mechanical properties through a high- temperature heat treatment of the

Nb 3 Sn, and low thermal expansion are required attributes of an alloy for

CICC use. Alloy composition has been adjusted to improve properties.

Fracture toughness, Kj c (J), is very high, 240 MPa-m^ at 4 K after aging.

Material that has been cold-worked, then aged can have yield strengths from

1489 to 1320 MPa, depending upon heat treatment. As noted above in the

discussion of the presentation by R.P. Reed on ITER fatigue problems,

fatigue crack growth rate in the short-crack region is an important

parameter for a CICC alloy. Fatigue crack growth rate tests reported here

were for constant AK (stress intensity range) , and did not cover the shortcrack region.

Cryogenic Properties of Al-Li Alloys : J.W. Morris, Jr., Center for Advanced

Materials, Lawrence Berkeley Laboratory, Berkeley, CA.

For many structural alloys, such as the austenitic stainless steels,

the increased yield strength at cryogenic temperatures is correlated with

some decrease in ductility and fracture toughness. However, because of a

change in fracture mode at lower temperatures, this effect does not occur

for the Al-Li alloy 2091, and fracture toughness at 4 K is higher than it is

at room temperature. Problems in welding Al-Li alloys and their use in

superplastic forming were also discussed.

Large Scale Cryogenic Test Facility at NIST : J.D. McColskey, R.P. Walsh, W.

Fehringer, National Institute of Standards and Technology, Boulder, CO.

Large scale cryogenic testing is required for the development of large

superconducting magnetic energy storage (SMES) systems. NIST has recently

acquired the capability to test materials under compressive fatigue loads up

to 4.4 MN (1 million lb) at temperatures from 4 to 300 K. Thermal isolation

members in a SMES system will experience such loading conditions and will be

tested by NIST. This is the only facility in the U.S. with this capability.

Details of the system were described.

STRUCTURAL ALLOY WELDING

Mechanical Properties of Low Ferrite Welds at 4 K : C.N. McCowan and T.A.

Siewert, National Institute of Standards and Technology, Boulder, CO.

The strength and fracture toughness of welds made with three

commercially available stainless steel electrodes (types 330, 316L, and 904)

were evaluated at 4 K. Nitrogen additions to the shielding gas were varied

to change the N content in the gas metal arc (GMA) welds. The type 904

welds were found to have the best combination of yield strength and

toughness at 4 K: 942 MPa and 390 and 269 MPa^m1* (2 tests).

351
Weld Alloy Design : T.A. Siewert, National Institute of Standards and

Technology, Boulder, CO.

A summary of progress at NIST during the past five years in

understanding the parameters that govern weld properties for cryogenic

applications was given. Critical parameters, such as N, Ni, and 0 content

in weldments, have been identified and quantified for some electrode-base

metal systems. Welds of AISI 304L and 21-6-9 failed in a brittle mode at

cryogenic temperatures. A transformation- stable electrode material,

designated 904, looks promising for cryogenic applications and may have

properties in the so-called JAERI-box region where yield strength >1200 MPa

and Kj c > 200 MPa*m^. (JAERI is the acronym of the Japan Atomic Energy

Research Institute and the JAERI box refers to a region of desired

properties on a Kj c vs. yield strength graph.) It was emphasized that no

single property can provide adequate characterization of a welded alloy. A

welded structure must be considered as a combination of the base metal, weld

metal, and the heat affected zone (HAZ) . Since these three zones are

interrelated, the study of a welded structure is complex.

Welding Problems. Panel Discussion : T.A. Siewert (Leader), National

Institute of Standards and Technology, Boulder, CO.

The possibility of using welded 21-6-9 stainless steel alloy in

cryogenic applications such as the CIT was discussed. Improved control of

metallurgical variables in the current production of this alloy may lead to

better properties than were reported in the past. Because of its very high

strength, 21-6-9 keeps resurfacing in new projects as a candidate alloy for

cryogenic applications. However, a considerable history of past

measurements indicates that cryogenic toughness, especially of welds, is

poor.

The problem of weld inspection was discussed. Designs of many fusion

systems will not permit much inspection of welds after construction is

complete. If AISI 316LN is used in thick- section welds for TF coils for

ITER, this will be a potential problem.

The composition (in wt.%) of the promising new electrode material for

cryogenic applications, 904, is Fe-25Ni-20Cr-5Mo. The compatibility of this

material with various base metals must be examined so that it can be used in

cryogenic applications.

METAL AND COMPOSITE TEST STANDARDS

Discussion: Low Temperature Mechanical Property Test Standards: Structural

Alloys and Composites : R.L. Tobler (Leader), National Institute of

Standards and Technology, Boulder, CO.

Standards for cryogenic measurement of both tensile properties and

fracture toughness have been progressing steadily through ASTM in recent

years. However, standards for composite testing, even at room temperature,

are not yet developed. Some new tests, such as the shear test based on the

352
Iosipescu configuration, are now under consideration for use at cryogenic

temperatures. Miniaturization and scaled down apparatus may be a practical

necessity for tests carried out in cryostats. The applicability of such

scaled-down measurements to actual design conditions then needs to be

considered.

The Effect of Strain Rate on Low Temperature. Round Robin Tensile Tests : R.P. Reed, R.P. Walsh, and R.L. Tobler, National Institute of Standards and

Technology, Boulder, CO.

The transition to discontinuous yielding in 4 K tensile testing occurs

at a lower strain rate for very high strength alloys (yield strength - 1400

MPa) . Since the measured ultimate tensile strength is a function of strain

rate during the test, lack of standards for the strain rate may explain the

different results found in recent round- robin testing.

Warm Prestress Effects in Cryogenic Structural Steels : R.L. Tobler, L.M.

Ma, and R.P. Reed, National Institute of Standards and Technology,

Boulder, CO

Preloads applied to steel specimens at a relatively high temperature,

Ti , can influence fracture toughness at a lower temperature, T2

. The effect

has been best known in ferritic steels, where T^ is above the ductile- tobrittle transition temperature, and T 2 is below it. However, this

presentation showed that the effect also occurred in some austenitic steels

when precracking at room temperature (T^) preceded fracture toughness

testing at 4 K (T2

). The effect increased the toughness measured at 4 K,

and was more important for steels of low toughness. It is apparently

associated with a change in the fracture mode between T^ and T2

. To

standardize test methods, the effect must be avoided in order to measure the

inherent toughness at 4 K. Japanese fracture toughness testing is not

comparable with U.S. testing in some cases because the Japanese generally

precrack specimens at room temperature.

J-Inteeral Test Analysis Procedures and Interpretation : P.T. Purtscher,

National Institute of Standards and Technology, Boulder, CO

A series of J- integral fracture tests were performed at 76 K on an Fe22Cr-13Ni-5Mn-2Mo-0. 3N steel. The data were analyzed in two ways: first,

with the initial ASTM standard (E 813-81) in which Jjc is derived from the

intersection of the blunting line and the tearing line; and second, with the

new standard (E 813-87) in which Jjc is determined from the crack growth

resistance at an offset of 0.2 mm. The new method was shown to produce

consistently higher values of Jjc with less scatter.

Negative crack growth was also discussed. Since the plastic zone is

not always ahead of the crack, but also at the sides of the crack,

compression may cause negative crack growth. The impracticality of Charpy

toughness tests below 77 K, and the lack of correlation of 77 K Charpy

results with 4-K Jjc toughness were also noted. Specimen warming at 4 K,

partly due to low specific heats, results in invalid Charpy tests.

353

INTERNATIONAL THERMONUCLEAR EXPERIMENTAL REACTOR

SPECIALISTS MEETING ON MAGNET MATERIALS, February 26-28, 1990

Introduction

ITER is a large international fusion energy project. Resident European,

American, Soviet, and Japanese specialists occupy two buildings of the Max

Planck Institute in Garching, near Munich, West Germany. On February 26-28,

resident and visiting representatives from participating countries met in

Garching to discuss cryogenic material requirements. The goal was to

formalize R&D plans for ITER magnets. The workshop was organized and

chaired by John Miller of Lawrence Livermore National Laboratory. NIST

representatives (R.P. Reed, N.J. Simon, and R.L. Tobler) contributed

technical presentations and drafted 5 -year plans for future research. The

format was presentations by the participants followed by preparation, and

discussion of research plans. Three sessions (each lasting one day) were

devoted to the material categories described below.

Structural Alloys and Welds

The structural alloys session focussed primarily on short crack fatigue

problems in the jacket of the superconductor cable- in- conduit configuration

for the toroidal field coils and on the production and properties of

austenitic steels in 400-imn thick sections of the toroidal field coil case.

Simon presented extensive work on parametric optimization of 316 LN alloys

and an evaluation of the 316 LN database at 4 K. In general, 316LN has

adequate properties for ITER applications, and the properties can be

tailored by control of grain size, purity, and composition. As a result , it

is clear that much multinational testing will be needed for the various

developmental alloys being advocated by Japan and the USSR if databases for

these alloys are to become as large as that for AISI 316 LN.

Simon also presented a brief summary of the alloy activation results

obtained at the TRIGA reactor of the Atominstitut in Vienna, through

Professor Weber. Samples of four steels were activated: (1) 316LN, (2) 310,

(3) 18Cr-18Mn, and (4) 25Mn-5Cr. A fast neutron fluence (E > 0.1 MeV) of 2

x 1022 m" 2 was used. This corresponds approximately to the fluence expected

at the magnet location from previous studies (MARS and STARFIRE) . The

lowest activity was obtained for sample (4), next lowest was sample (3),

then sample (2), and finally sample (1) 316LN had the highest activation.

Low Cr and Co contents were important in determining this ranking. The

TRIGA spectrum does not contain 14 Mev neutrons. When a theoretical

calculation was carried out for activation of samples (1) and (4), under a

neutron fluence of 1 x 10 22 m' 2 at 14 MeV, Mn and Fe contents dominated and

the activation position of samples (1) and (4) reversed at a time

immediately after irradiation.

355
Tobler reported mechanical property measurements for a candidate ITER

conduit steel. These data for JK1 (a niobium-modified 316 LN steel) were

obtained thanks to assistance from J. Berger, R.P. Walsh, R. Santoya, and C.

McCowan. NIST provided ITER designers with the first short -crack simulation

fatigue data ever measured at 4 K. These fatigue data are essential to safe

life predictions. However, we also reported low toughness for this

material; the toughness values in terms of Kj c

(J) were about half of that

reported in developmental studies by the steel manufacturers. This large

disparity showed the importance of material specifications, test standards,

and verification testing, since fracture toughness minimums were not

specified in the steel procurement. Although the tests for JK1 gave the

impression that the hoped-for improvement in properties of the conductor

jacket might be delayed, or more difficult to realize, alternative alloys

will be evaluated at NIST.

Reed contributed a talk on structural alloys and later was asked to write

the 5-year plan for structural alloys. These plans include recommendations

on screening of candidate alloys for case and conduit applications, welding,

design allowables, and development of appropriate low temperature data

bases. Considerable research is required in view of the plan to use very

thick (400 mm) austenitic steels for the magnet case structures. Both Reed

and Summers (LLNL) articulated the importance of production problems for the

very thick sections required. The total tonnage required by ITER (50 000)

is insufficient to warrant special considerations from mill facilities. A

possible alternative, however, is to consider construction of thick sections

from thinner sections.

Other items

:

-- Muster and coworkers (EMPA, Zurich) report that in 316 LN type alloys

only carbides form during sensitization, not nitrides. We can

eliminate sensitization effects in nitrogen- strengthened austenitic

alloys (if EMPA is correct) by reducing the carbon content to about

0.01. This would constitute a metallurgical breakthrough for

austenitic steels.

Nyilas (KfK, Karlsruhe) presented data for low- temperature serrated

yielding in austenitic steels and his results are identical to those

reported by NIST last year.

The Efemov Institute (Leningrad) is now collaborating with Yushchenko

(Paton Institute, Kiev) in structural alloy development. Alekseev of

Efemov also encouraged collaboration between cryogenic structural

material groups in the US and USSR. (Yushchenko first broached this

subject at the International Cryogenic Materials Conference in 1989.)

Superconductors

Nb3 Sn superconductor cable development was discussed during the second day's

session. Nb3 Sn is needed because the ITER magnets must produce fields of

356
about 12 T with high current density. Compared to NbTi, the NbjSn conductor

is in its infancy. As Gregory of IGC explained, there is little or no

market at present. Nevertheless, IGC appears to have developed a product

that has very favorable J c and Tc properties. Miller asked participants to

respond to questions regarding the size of billets required for development

of previous superconductor stabilizer composites, the length of conductor

that could be produced, and so forth, to estimate the amount of additional

development that would be required to realize Nb 3 Sn conductors.

Magnet Insulation

The third session concerned needs for magnet insulation (composites) . The

Soviets described their cryogenic irradiation facility. Reed reported

NIST's plans to develop cryogenic shear tests for bulk and miniaturized

specimens. One proposed use of nonmetallics is in very thin layers (about 2

mm) between ITER magnet windings; bulk material tests are inappropriate for

this application, and the thin layer shear measurement methods (lap shear)

developed by the Europeans have not yet received adequate international

validation. The method of construction for ITER magnets is expected to

involve vacuum impregnation with an epoxy resin of very long pot life (at

least 24 h at room temperature) . Other parts of the proposed magnet

structure require 12- to 25-mm plates of glass-reinforced epoxy.

Previously, NIST contributed to the development of G—10CR, a standard

cryogenic grade material that may be applicable for these thicker section

structural parts. Bulk test methods are useful for such applications, and

they may provide a scaled-up method for validating the lap shear tests

needed to evaluate the impregnated insulation after irradiation.

Reed was asked to write a 5-year plan for ITER research on insulators. This

plan emphasizes development of radiation resistant insulation. Filament

configuration and polymer matrix should be optimized for resistance to shear

forces at 4 K, after irradiation at 4 K.

ITER R&D STRUCTURAL MATERIALS PROGRAM

I. ALLOYS

The current design requirements of ITER for structural alloys for the magnet

case and cable conduit exceed the existing technology base of use at

cryogenic temperatures. Very thick sections (-400 mm) of high-strength

(>1000 MPa yield strength) austenitic steels need to be welded and inspected

to use as the magnet coil case. The cable conduit requires high strength,

excellent fatigue crack growth resistance at 4 K, and the ability to

withstand, without metallurgical deterioration, the superconducting

processing schedule of a long aging period at about 700°C. These

requirements, for structural alloys usage at 4 K, are new and necessitate a

sizeable R&D effort. This R&D effort is outlined here.

357
Phase I: (*”1 year) : Down select to 2 alloys for case structure and 1 or 2

alloys for conduit. Principal criteria for case alloy selection

are strength and toughness; for conduit alloy, strength and

fatigue -crack growth rate. Screening tests on these properties

for candidate alloys should be run at several national labs.

Phase II: Case Materials

Two classes of alloys should be tentatively included for Phase II study:

(1) 316LN, an alloy class that has been extensively studied for the

strengthening effects of nitrogen, grain size, and temperature and for which

high speed weld techniques are being developed. This alloy should be

selected if yield strengths of about 1000 MPa, coupled with a fracture

toughness of about 250 MPa-m^, are required at 4 K; (2) an alloy class that

has yield strengths of about 1200 MPa and fracture toughness of about 200

MPa*m^. Candidate alloys include JN1 (Japan); 2QCr-6Ni-6Mn (USSR);

13Cr-9Ni- lOMn (USSR).

Special considerations for case alloys are the requirements of very thick

sections (-400 mm) . This thickness has never been used before in a critical

structure at cryogenic temperatures. Research studies (Phase II) include:

Development of thick- section processing techniques (rolling and/or

forging)

.

Development of thick-section welding techniques and filler rod for

high-strength austenitic steels.

Development of inspection capabilities (NDE)

.

Scale-up studies from specimen coupons to nearly full-size

thickness and geometry. Measurements include effects of thickness

(reflecting variability from grain size and orientation) on

strength, toughness and fatigue. Large test facilities (> 500

tons) at 4 K are required.

Compilation, critical review, and measurement of those physical

and mechanical properties required in design.

Development of specifications for chemistry and control, allowable

grain size and thermal/mechanical processing.

Establishment of allowable flaw sizes of weldments.

Establishment of allowable stresses.

Phase II: Conduit Materials

The structural alloy for the conduit must have low fatigue-crack growth

rates with high yield strength at 4 K. A number of candidate alloys exist,

but some have not fulfilled expectations. These alloys include JBK-75 and

Incoloy 908 (U.S.), JK1 (Japan), and 316LN with very low carbon (A410,

Europe)

.

358
Research studies (Phase II) include:

Measurements of fatigue crack growth rate, including long and

short-crack simulations, on base alloys and welds.

Development of welding techniques.

Development of inspection technique.

Establishment of allowable flaw sizes.

Establishment of property variability leading to

establishment of allowable operating stress and

maximum fatigue cycles.

Study of effects of spectrum loading on fatigue

crack growth rates at 4 K.

Full-size simulation test of conduit, in fatigue and

tensile load.

Compilation, critical review, and measurement of

those physical and mechanical properties required

in design.

II . INSULATORS

In current ITER design, only thin layers of potted epoxy base, with tape

reinforcement, will be used. Projected layer widths are 0.03-0.05 mm.

Thus, for this design that has an absence of bulk insulators, only research

to optimize thin- sheet, reinforced epoxy insulation should be supported.

This includes the following items:

Development of irradiation- resistant (5x10^ Gy) epoxy with a long

(-24 h) pot life and low viscosity at room temperature. (Most

epoxies that are currently used do not have a long pot life at

room temperature .

)

Development and correlation of shear/compression test techniques

for comparing thin- sheet materials and the effects of irradiation

at low temperatures. (Appropriate test techniques are recommended

in the section under Test Techniques for Structural Insulators.)

Irradiation studies with a higher ratio of neutron to gamma

radiation and including more high-energy neutrons are suggested.

Parallel irradiation of specimens with gamma radiation only is

necessary to separate out the damage from neutrons and gamma rays

.

Irradiation at 4 K, coupled with in situ tests at 4 K, are needed.

Irradiation at 4 K, followed by warming to 295 K and subsequent

testing at 4 K are also needed to study Wigner (defect

accumulation) and gas accumulation effects.

359
Dielectric breakdown tests as a function of compression/shear

loading should be conducted at 4 K.

Compression :

Specimen width- to-height ratios must be large enough to prevent column or

shell buckling. Constraint (support) at specimen ends eliminates brooming

effects and leads to increased strengths. Otherwise, compression testing is

relatively simple.

Shear

:

There is a wide diversity of shear test methods and test results at

cryogenic temperatures. For continuous, filamentary- reinforced composites

in bulk form (plate, rod), these include:

torsion

tension (± 45°, 10° off axis)

rail shear

short-beam shear

4-point flexure

notched (Guillotine) shear

double -lap shear

Iosipescu test

AFPB*

For thin sheet composites, the following test techniques are used: singlelap shear and double -lap shear

A miniaturization of the single- lap shear test has been developed and used

for NET irradiation studies by Fardi, Mitchell, and Poehlchen. Finite

element analysis and preliminary test results indicate that the test

techniques may be an acceptable alternative for small specimens

.

Of special interest is the new technique developed by Okada's group that

permits the simultaneous application of both shear and compressive loads.

Results from these tests suggest that the shear strength is strongly

dependent on simultaneously applied compression loads, increasing with

higher load.

For ITER insulation, only very thin sheets of filament -reinforced composites

are projected at present. Thus, only single or double lap shear, or the

shear/compression type tests developed by Okada will be useful for the

selection of a long pot- life epoxy and to measure the shear strength of the

selected insulator.

For testing of unirradiated materials, it is recommended that a new test be

developed for use as a "standardized ITER shear test." For testing of

irradiated insulators, Fardi et al.'s miniaturized lap shear test should be

developed and correlated with the newly developed test.

360
APPENDIX
.
MECHANICAL PROPERTIES OF ULTRA-LOW CARBON STEEL

AND EFFECTS OF WARM PRESTRESS AT 4 1C*

R.L. Tobler, L.M. Ma*, and R.P. Reed

Materials Reliability Division

National Institute of Standards and Technology

Boulder CO 80303

The yield strength of ultra- low carbon steel increases from

137 MPa at 295 K to 705 MPa at 4 K. In liquid helium at 4 K the

material is brittle and cleaves in tension. Consequently it is

difficult or impossible to properly fatigue -precrack toughness

specimens for tests at 4 K. In this study, we precracked a series

of compact specimens at room temperature (295 K) and then fractured

them at 4 K. The toughness measurements at 4 K are affected by

precracking at 295 K, and they provide only an estimate of Kj c for

this steel which is 15 MPa-m^5. The tensile properties of the steel

at 295 and 4 K are reported, and the effects of prestress on the

tensile and fracture properties at 4 K are discussed.

INTRODUCTION

Owing to their ferritic microstructures and lack of nickel content,

carbon steels undergo ductile -to -brittle transitions at temperatures ranging

from 250 to 375 K [1-5]. The transition temperature marks an abrupt change

in fracture resistance and it determines the minimum service temperature. On

the upper shelf, ductile fracture occurs by microvoid growth and coalescence;

on the lower shelf, cleavage fracture occurs with little or no ductility.

Normally, this transitional behavior of carbon steels disqualifies their

use as primary load carrying members in cryogenic structures subjected to

tensile stress. In an exceptional circumstance, ultra- low carbon steel at a

subtransition temperature of 4 K is being considered for magnet yokes in the

Superconducting Super Collider. No other steel has the desired ferromagnetic

properties for this application. The proposed application prompts serious

questions about the maximum load carrying capability of the carbon steel at 4

K, and to what extent it might be affected by prestressing (preloading, or

prestraining) at room temperature. The measurements described here provide

clarification.

* Guest worker at NIST, on leave from Institute of Metal Research, Academia

Sinica, Shenyang, P. R. China.

363
Discussions [6, 7] and reviews [8, 9] of warm prestress (WPS) described

how preloads applied to steel specimens or structures at a relatively high

temperature T can increase the toughness subsequently measured at a lower

temperature, T . Toughening from WPS is common in carbon steels when T is

above the transition temperature and T is below the transition temperature.

Toughening was reported in notched or fatigue precracked specimens of high-C

spheroidized steel [6], medium strength steels [10], C-Mn steels [11], 0.36%

C steel [12], A 533 B-l steel [13], and free machining steel [14]. In this

paper, we describe the effects of WPS for an ultra- low carbon steel where T

is room temperature and T is liquid helium temperature (4 K) .

1

MATERIAL AND SPECIMENS

The test material was a commercial ultra- low carbon steel plate, 89 mm

thick, with the following composition (mass percent): Fe-0 . 007C-0 .021Mn0 . 012P-0 . 014S -0 . 006Si -0 . 052A1-0 . 003N . Compared to conventional mild carbon

steels, this steel has a lower C and Mn content which confers the desired

ferromagnetic behavior at cryogenic temperatures.

Six tension test specimens in the short- transverse orientation were

machined from the as -received plate and tested following ASTM Methods

E 8-89 [15]. Five compact specimens in the S-L orientation were also

machined and tested (after WPS) following ASTM Methods E 399-83 [16].

Figure 1 shows the test specimen orientations with respect to the 89 mm

thick plate. Tension specimens were cylindrical with a diameter of 6.35 mm,

a length of reduced section of 38.1 mm, and a gage length of 25.4 mm.

Displacement was measured using clip-on extensometers . The specimen reduced

section was tapered slightly to ensure fracture at the center of the gage

length. Fracture toughness specimens were 25.4 mm thick (W) and 50.8 mm wide

(W)

,

with the planar proportions recommended by Methods E 399-83.

PROCEDURE

The apparatus for conducting cryogenic tension and toughness tests were

described before [17-19]. The test procedures for measuring the tensile

properties and stress intensity factors for compact specimens are specified

in ASTM Standard Test Methods E 8-89 and E 399-83, respectively.

Tension tests

The crosshead velocity during tension tests was 0.5 mnrmin" 1 which

corresponds to a nominal strain rate of 0.013 min" 1

. Two tests were

performed at 295 K, and two at 4 K. Two more tests were conducted to observe

the effect of prestrain at 295 K on tensile properties at 4 K; in these tests

the specimens were prestrained 5% at 295 K, then unloaded, cooled in liquid

helium, and fractured at 4 K.

364
Figure 1. Test specimen geometries and orientations.

Fracture Toughness

Valid Kj c data could not be measured due to difficulties encountered in

fatigue -precracking. Usually, test specimens are fatigued in tension and Kj c is then measured at the same temperature. Precracking in tension is routine

at temperatures where the material retains ductility (see paragraph A2.1.1 of

Method E 399-83), but it is difficult or impossible to fatigue brittle metals

in tension at extreme cryogenic temperatures [20, 21]. We therefore

precracked the specimens at 295 K where the steel is ductile, although this

affects the toughness at 4 K.

Five specimens were fatigued at 295 K, using a minimum- to -maximum load

ratio of 0.1. The waveform was sinusoidal in load control at 25 Hz, and

crack length was monitored by compliance [19]. The final maximum fatigue

stress intensity factor, K^x, varied from 11 to 50 MPa-m

1*, as Table 2 shows.

365
The two lowest Kq^, 11.7 and 16.5 MPa'm**, were achieved by long* term cycling

with periodic load reductions as a function of crack growth. After reaching

a crack- length- to -width ratio of ~0.5, the specimens were fractured at 4 K at

a displacement rate of approximately 0.03 mms" 1

.

The above procedure is a fatigue- load, unload, cool, and re- load- to

fracture sequence which toughens the steel due to WPS. Since the apparent

toughness at 4 K increases with the magnitude of room- temperature fatigue

loads, we applied progressively lower K^x to diminish the effect. We

plotted the apparent 4-K toughness versus the plastic zone size associated

with the final maximum fatigue load at 295 K, then extrapolated the trend to

estimate Kj c

. The plastic zone radius Rp was estimated from the yield

strength YS and K^x at the appropriate temperature using Irwin's

approximation [22]:

Rp - (l/6ir)*(KmaxAS) 2

(1)

RESULTS

Tensile properties

The tensile mechanical property measurements are listed in Table 1. The

data for specimens numbered 1—4 establish the baseline tensile properties for

the ultra- low carbon steel. The average strength increases substantially at

4 K while ductility falls nearly to zero. Between 295 and 4 K the yield

strength shows a fivefold increase, the ultimate strength nearly triples, the

elongation drops from 48.7 to 1.2%, and the reduction of area drops from 89.4

to 0.5%.

Table 1. Tensile properties of ultra-low carbon steel at 295 and 4 K.

Specimen

No.

Temperature

(K)

Yield

Strength

(MPa)

Ultimate

Strength

(MPa)

Elongation

(%)

Red. of

Area

(%)

1 295 137 267 49.4 89.4

2 295 137 269 47.9 89.4

Average 137 268 48.7 89.4

3 4 700 753 1.4 0.5

4 4 709 735 1.9 0.5

Average 705 744 1.2 0.5

5 295, then 4

a NA 679 5 .

3

b 4 .

0

b

6 295, then 4

a NA 702 5.0b 4 .

8

b

Average NA 690 5 .

2

b 4.4b

^restrained approximately 5% at 295 K, unloaded, cooled to 4 K, and

then re -loaded to fracture at 4 K.

bthese values reflect ductility during the initial loading at 295 K.

366
Table 1 also lists data for specimens (numbers 5 and 6) that were warm

prestrained at 295 K before testing at 4 K. The 295-K prestrain slightly

lowers the fracture stress of unnotched specimens at 4 K. The prestrained

specimens failed at an average stress of 690 MPa; this is nearly equal to the

yield strength of specimens that were not warm prestrained, and is about 7%

lower than the ultimate strength of specimens that were not prestrained.

Fracture toughness

Load—displacement records for compact specimens are shown in Fig. 2.

Most tests satisfy Method E 399-83, except the fatigue load restrictions

(Section A2.4.4) were violated by cycling at 295 K to create WPS. Also, the

Pm/PQ rati- 0 for specimen MS-6 slightly exceeded 1.10 which is the allowable

limit specified by the standard (Section 9.1.2).

Table 2 lists the toughness calculations as Kq, not Kj c , because they

are invalidated by warm precracking. Kq decreases systematically from 36 to

23 MPa-m^ as Kmax is lowered, but, as shown in Fig. 3, no plateau is reached

at lower K^x- In other words, the effect of WPS was reduced but not

eliminated by at the low fatigue loads used in this study. We conclude that

the lowest measured Kq value, 23 MPa-m^, is still affected to a degree by

WPS, and therefore it represents only an upper bound for Kjc at 4 K.

With some assumptions these data can be used to estimate Kj c

. A plot of

Kq (at 4 K) versus plastic zone size approximations corresponding to K^x (at

295 K) is shown in Fig. 4. The 4-K toughness varies semilogarithmically with

the estimated plastic zone size at the crack tip. The estimated plastic

zones after fatigue at 295 K are large (6.9 to 0.39 mm), because the steel's

yield strength is relatively low at 295 K. A much smaller deformed zone

would form if precracking could be accomplished at low temperatures where the

resistance to deformation is higher. Using the upper bound value of

toughness (23 MPa'm^) and the 4-K yield strength (705 MPa) in Eq. (1), we

estimate the plastic zone size for a valid 4 K test to be 0.05 mm. An

extrapolation of the trendline in Fig. 4 to 0.05 mm then gives a Kjc estimate

of 15 MPa’m14 for this steel with no WPS at 4 K.

Table 2. Precracking conditions and 4-K fracture toughness results.

Specimen

No.

Maximum Stress Intensity

Factor Used In Fatigue

at 295 K, K^x

(MPa -in

1 )

Fracture Toughness

Measured at 4 K, Kq

(MPa-m*5

)

MS-3 49.4 35.6

MS-5 32.9 32.6

MS-2 22.0 28.3

MS -4 16.5 26.5

MS-6 11.7 23.

2

a

aInvalid: the ratio Pm

/PQ is 1.12; a ratio of 1.10 or less

is permissible (Method E 399-83, Section 9.1.2).

367
368

Figure

2.

Load-displacement

curves

for

mild

steel

compact

specimens

at

4

K
z/tUiedW^ Me 0*

369

Figure

3.

Fracture

toughness

Kq,

at

4

K,

as

a

function

of

K

raax

applied

previously

at

295

K.
mw.ediAi ‘°>i ‘ssaNHonoi aaniovaa

o

XI

>0 a

a> 10 o M

o 4J u x

cu a> <

<yM

3

00

•H

!X

370

PLASTIC

ZONE

RADIUS,

R

D,

mm
DISCUSSION

The effect of warm prestrain on unnotched specimens is to lower the

fracture stress slightly at 4 K. Prestrain at room temperature further

compromises the steel's ability to deform on reloading at 4 K. So, unlike

the as -received steel, the 295 -K prestrained material could not sustain a

0.2% plastic strain at 4 K (see Fig. 5) and failure in the prestrained

specimens occurred at slightly lower tensile stresses.

Contrary to the effect on smooth tensile specimens, WPS is beneficial

for deeply cracked compact specimens. Reasons for the toughening of notched

or cracked specimens are discussed in the literature [6—9], In some

practical structural applications with linear-elastic materials, the

beneficial effect of WPS is used in a strategy of intentional overstressing

at T to reduce the probability of fracture at [8]. In proposed magnet

applications, however, it is not clear how WPS could be practically used.

Precracking at low loads is time consuming. In principle, a valid Kj c measurement at 4 K is approached as the precracking loads at 295 K are

lowered. However, the time required for precracking this steel is

prohibitive at the required low loads. We therefore resorted to estimating

the toughness.

CONCLUSIONS

1. The ultimate strength of ultra- low carbon steel at 4 K is 744 MPa. Room

temperature prestraining in tension by 5% lowers the ultimate strength

at 4 K to 690 MPa.

2 . The fracture toughness (Kj c ) of this ultra- low carbon steel at 4 K is

estimated at 15 MPa-m

1*. Valid Kj c measurements could not be obtained by

fatigue precracking in tension at a stress ratio of 0.1. Fatigue

precracking at 295 K inflated the toughness measurements at 4 K, whereas

fatigue precracking at 4 K was impossible because of brittleness.

ACKNOWLEDGMENT

The test material and funds for conducting this work were provided by

Brookhaven National Laboratory, with M. Rehak as project monitor.

REFERENCES

[1] Lubahn, J.D., Effect of Temperature on the Fracturing Behavior of Mild

Steel, Welding Journal Research Supplement, Vol. 35, 1956, pp. 557S568S

.

[2] Wessel, E.T.

,

A Tensile Study of the Brittle Behavior of a Rimmed

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371
800

in

Mi

<

CC

H- CO

(

m

<Du

3

60

•r-l 1*4

BdlAI

* ‘SS3U1S

372

Stress

versus

strain

curve

for

a

mild

steel

specimen

(A)

prestrained

5»

at

295

K,

then

(B)

reloaded

to

failure

at 4

K.
[3] Hahn, G.T.

,

Averback, B.L., Owen, W.S., and Cohen, M. , Initiation of

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Wiley, New York, 1959, pp. 91-116.

[4] Clausing, P. , Tensile Properties of Eight Constructional Steels

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473-492.

[5] Shoemaker, A.K. and Rolfe, S.T., The Static and Dynamic Low

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[7] Curry, D.A.

,

A Model for Predicting the Influence of Warm Prestress

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[8] Nichols, R.W.

,

The Use of Overstressing Techniques to Reduce the Risk

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Notch Fracture Strength, Journal of Basic Engineering, Vol. 85, 1963,

pp. 97-101.

[11] Harrison, T.C., and Feamehough, G.D., The Influence of Warm

Prestressing on the Brittle Fracture of Structures Containing Sharp

Defects, Journal of Basic Engineering, Vol. 94, 1972, pp. 373-376.

[12] Nichols, R.W.

,

Levine, W.H.

,

Quirk, A., and Bevitt, E.A.

,

A Limit

Approach to the Prevention of Pressure Vessel Failure, Proceedings of

the First International Conference on Fracture, Sendai, Japan, 1966,

pp. 1673-1678.

[13] Nakamura, H. , Kobayashi, H. , Kodaira, T. , and Nakazawa, H. , On the

Effects of Pre-loading on the Fracture Toughness of A533 B-l steel,

in: Advances in Fracture Research, Proc. ICF 5, Vol. 2, Pergamon, New

York, 1981, pp. 817-824.

[14] Curry, D.A.

,

A Micromechanistic Approach to the Warm Prestressing of

Ferritic Steels, International Journal of Fracture, Vol. 17, 1981, pp.

335-343.

373
[15] Standard Test Methods of Tension Testing of Metallic Materials, ASTM

Designation: E 8-89, in: 1989 Annual Book of ASTM Standards, Section

3, Metals Test Methods and Analytical Procedures, Vol. 03.01, Amer.

Soc. Test. Maters., Philadelphia, 1989, pp. 131-146.

[16] Standard Test Method for Plane Strain Fracture Toughness of Metallic

materials, ASTM Designation: E 399-83, in: 1989 Annual Book of ASTM

Standards, Section 3, Metals Test Methods and Analytical Procedures,

Vol. 03.01, Amer. Soc. Test. Maters., Philadelphia, 1989, pp. 487-511

[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, JTEVA, Vol. 13, No. 4, 1985, pp. 299-302.

[18] Reed, R.P., A Cryostat For Tensile Tests in the Temperature Range 300

to 4 K, in: Advances in Cryogenic Engineering, Vol. 7, Plenum, New

York, 1961, pp. 448-454.

[19] Read, D.T.

,

and Tobler, R.L.

,

Mechanical Property Measurements at Low

Temperatures, in: Advances in Cryogenic Engineering, Vol. 28, Reed,

R.P. and Clark, A.F., Eds., Plenum, New York, 1982, pp. 17-28.

[20] McCammon, R.D.

,

and Rosenberg, H.M.

,

The Fatigue and Ultimate Tensile

Strengths of Metals Between 4.2 and 293 K, Proceedings of the Royal

Society of London, Vol. A 242, 1957, pp. 203-211.

[21] Stonesifer, F.R.

,

Effect of Grain Size and Temperature on Fatigue

Crack Propagation In A 533 B Steel, Engineering Fracture Mechanics,

Vol. 10, 1978, pp. 305-314.

[22] Irwin, G.R.

,

Plastic Zone Near a Crack and Fracture Toughness, in:

Mechanical and Metallurgical Behavior of Sheet Materials, Proc. 7th

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Ordnance Research, 1960, pp. 63-77.

374
CRYOGENIC FATIGUE OF HIGH-STRENGTH ALUMINUM ALLOYS

AND CORRELATIONS WITH TENSILE PROPERTIES 4

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

Fracture and Deformation Division

National Institute of Standards and Technology

Boulder, Colorado 80303

ABSTRACT

Notched and unnotched sheet specimens of four aluminum alloys were

fatigue tested to evaluate potential use in the Superconducting Super

Collider. Alloys 7075-T6, 7475-T761, 2219-T87, and 2090-T8E41 were tested

in axial fatigue at a stress ratio of 0.1. The unnotched specimens were tested at 295, 76, and 4 K, whereas notched specimens were tested at 76 K

only. Ue compare the fatigue strengths of these alloys with a practical

interest in a life of 10 6 cycles. Ue also correlate fatigue strengths with

static tensile strengths for notched and unnotched specimens.

INTRODUCTION

High-strength aluminum alloys have acquired considerable importance

because of their current and potential uses in aerospace structures and

superconductor projects. 1 Superconducting Super Collider (SSC) dipole

magnets, for example, will probably require an aluminum alloy with high

fatigue resistance at low temperatures. 2 Since conventional fatigue data

are not readily available in the open literature, we report smooth and

notched specimen data in this paper to aid material selection. Strong

temperature effects on the fatigue lives of these aluminum alloys are demonstrated, and the possibility of predicting the fatigue resistances of

these alloys on the basis of simple correlations with tensile properties is

examined.

MATERIALS AND PROCEDURE

Four commercial aluminum alloys were tested in their as—received

tempers: 7075—T6, 7475—T761, 2090-T8E41, and 2219—T87. Specimens were taken from sheets in the recrystallized condition that were about 1.6 mm

thick. The nominal alloy compositions are:

+ Contribution of NIST, but not subject to copyright in the United States.

* Guest Researcher at NIST, on leave from the Institute of Metal Research,

Academia Sinica, Shenyang, P. R. China.

**Research Institute of Science and Technology, Pohang, South Korea.

Advances in Cryogenic Engineering (Materials). Vol. 36 1143
7075: Ai-5.6Zn-2.5Mg-l.6Cu-0.23Cr

7475: Ai-5.7Zn-2.3Mg-l.5Cu-0.22Cr

2090: Ai-2.86Cu-2.05Li-0.12Zr

2219 : Ai-6 . 3Cu-0 . 3Mn-0 . 18Zr-0 . 10V-0 . 06Ti

.

Figure 1 shows the specimen geometries used. All specimens were fabricated in the transverse orientation relative to the principal rolling

direction of the sheets. The square notch, shown in Fig. 1, essentially

duplicates the notched condition of an SSC key collar component 2 and has a stress concentration factor Kt of about 3.9. The specimens were ground on each side to remove surface marks from the as-received sheets. Unnotched

fatigue specimens were ground to number 1000 grit. Notched specimens were ground and then electropolished in a solution of 30% nitric acid and 70%

methanol. The specimen dimensions and notch root radii were measured using

a micrometer and a shadowgraph. The radii measurements at both corners and on both sides of each specimen were averaged. The minimum and mean values

of the four readings for each specimen were 0.555 mm ± 7% and 0.600 ± 10%,

respectively, for all alloys.

The specimens were mounted in a cryostat 3 filled with liquid nitrogen

or liquid helium for cryogenic testing. Tension tests were performed at a crosshead rate of 0.5 mm/min. The gage length for elongation measurements

was 25.4 mm. Tension- tension fatigue tests were conducted using a servo- hydraulic machine in the load-control mode. The minimum- to -maximum fatigue

stress ratio (R) was 0.1. To begin a fatigue test, the mean load was applied at once, and the load amplitude was adjusted while the frequency was gradually increased from 1 to 9 , 10, or 15 Hz (notched specimens), or to 15 or 20 Hz (smooth specimens).

RESULTS AND DISCUSSION

Smooth and Notched Specimen Tensile Properties

The conventional tensile properties and notched specimen tensile

strength (NTS) measurements are listed in Table 1. For all four alloys,

the yield strengths at 0.2% plastic strain (YS) and the ultimate tensile

strengths (UTS) increase with decreasing temperature between 295 and 4 K

while tensile elongations decrease moderately. Alloy 7075—T6 has the

highest strengths and 2219-T87 the lowest strengths at temperatures in this

range. The elongations are slightly temperature dependent, ranging from 9.9

to 17.1% (7475-T761 has the highest values).

Figure 1. Specimen geometries.

A. Smooth tensile specimen

B. Smooth fatigue specimen

C. Notched tensile or notched

fatigue specimen

T Note: All dimensions in mm.

r

ooi itdii. 0.31

376
Table 1. Static Tensile Properties for Smooth and Notched Specimens

data are mean values from two tests)

Alloy Temp. YS UTS El E NTS NSR NYR

K MPa MPa « GPa MPa NTS/UTS NTS/YS

7075 295 502 589 16.9 63.2

-T6 76 589 714 15.8 73.2 450 0.63 0.76

4 648 810 10.1 74.4

7475 295 460 515 17.1 66.3

-T761 76 549 636 17.3 75.3 472 0.74 0.86

4 572 739 15.1 76.4

2219 295 397 475 12.6 67.8 —T87 76 484 597 13.5 76.5 435 0.73 0.90

4 539 711 12.4 77.9

2090 295 488 528 12.1 74.0

-T8E41 76 551 640 10.8 76.9 443 0.69 0.80

4 614 727 9.9 84.1

Despite moderate ductility in the unnotched tension tests, the notchstrength ratios (NSR - NTS/UTS) and notch-yield ratios (NYR - NTS/YS) for

the alloys at 76 K range from 0.63 to 0.90, and are less than one, indicating notch sensitivity in each case. The NSR and NYR ratios are qualitative indicators of notch sensitivity and fracture toughness. In particular,

the NYR for round bar specimens correlates well with the plane-strain fracture toughness, Kjc

. 4 As previously illustrated by plotting YS versus NYR

for these alloys, the cryogenic toughness of our recrystallized 1.6 mm 2090-

T8E41 sheet alloy is not at all exceptional compared to the other alloys.®

Fatigue Life

The fatigue life data are plotted in Fig. 2 where Sm is the maximum

fatigue stress applied in each test, and Nf is the number of cycles to

failure. The studied range of Nf is between 10 4 and 10 7. For all alloys,

Nf increases with decreasing Sm, and the fatigue life curves improve by

substantial margins at cryogenic temperatures. Figure 3 compares the alloys

at 295, 76, and 4 K. Alloy 7075—T6 has superior fatigue resistance at all . three temperatures, and the ranking of the other alloys varies for specific

conditions as shown in Fig. 3.

Temperature Dependence of Fatigue Strength

Fatigue strength (FS) , defined here as the maximum fatigue stress cor- responding to a specified life, is plotted in Fig. 4 for a life of 10®

cycles (unnotched specimens). Vith decreasing temperature, FS increases

strongly, but the effect diminishes as Nf increases. At any temperature,

alloy 7075—T6 has the highest FS for Nf - 10® ; the other three alloys show

little difference for Nf - 10®, but there are differences at Nf - 10®, as can be seen from Fig. 3.

Reasons for the strong temperature dependence of fatigue resistance of

aluminum alloys were discussed previously.® The superior fatigue resistance

of 7075-T6 in the unnotched condition arises primarily from its higher

static ultimate strength. The materials have only modest tensile ductilities , and no obvious plastic deformations were observed on the fatigue

fracture surfaces of failed specimens in this study.

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377
Figure 2. Fatigue life curves for high strength aluminum alloys.

Since static tensile strength is quite important, we compare the alloys

on a normalized basis, using the ratio FS/UTS as a relative measure of

fatigue resistance. Figure 5 shows that the relative fatigue resistance

always increases with decreasing temperature. Whereas alloy 7075-T6 has the

highest FS at all temperatures, it does not have the highest FS/UTS ratio at

all temperatures; alloy 2090-T8E41 (not 7075-T6) has the highest relative

fatigue resistance at 4 K (FS/UTS - 0.95, for a life of 10 s cycles).

Notch Effect on Fatigue

Figure 6 shows the effect of the notch on S m -Nf curves. Notching

drastically decreases the fatigue lives of these alloys, but alloy 2090-

T8E41 emerges with the highest notched specimen fatigue strength (NFS) for a

Nf. Cycles

1146

378
•00

Figure 4. Temperature dependence

of FS, at Nf - 10 5

. Figure 5. Fatigue strength, at

Nf - 10* , normalized

by UTS.

life of 10 s cycles. The fatigue notch sensitivity factor (q) can be used to

rank the relative notch sensitivity of the alloys in fatigue:

( 1 )

Here, Kf - FS/NFS and Kt is the stress concentration factor for the notch.

When q - 0, it means that the fatigue life of the material is entirely

insensitive to the notch (FS - NFS) . On the other hand, the fatigue life is

notch sensitive to the degree that q approaches or exceeds 1. Table 2 lists our experimentally determined q and Kf values for comparison at 76 K and 10 s cycles. Of all four alloys, 7075—T6 is most sensitive to the notch, and

2090-T8E41 is least sensitive. Thus, 2090-T8E41 is superior to the three

other alloys on the basis of both absolute and relative notched specimen

fatigue resistance.

Figure 6. S-N curves of smooth and notched fatigue specimens for the four

aluminum alloys at 76 K.

1147

379
Table 2. Notch effects In terns of q and Kf (T - 76 K)

Alloy q Kf*

7075-T6 1.84 6.33

2219-T87 1.48 5.30

7475-T761 1.41 5.08

2090-T8E41 1 . 14 4.31

*Kf - FS (at 10 s

) /NFS (at 10 s

)

Fatigue Strength and Tensile Property Correlations

Smooth specimens. In Fig. 7, FS is plotted versus UTS for the four alloys

using all data at 295, 76, and 4 K. With some scatter, we find a clear

linear relationship between FS and UTS for the alloys at Nf - 10 s

:

FS ( 10 s

) - 1 . 42 (UTS) - 409 (MPa) (2)

This expression derives from a least squares fit, has a correlation coefficient of 0.96, and is rather precise; considering all the data including

T - 295, 76, and 4 K, the predicted and measured FS values agree to ± 22%

the agreement for 76 and 4 K data is ± 10%. Similar expressions for other

lifetimes can be derived, such as that shown in Fig. 7 for 10 6 cycles. The

capability of approximately predicting FS at 10 s or 10® cycles from UTS

values, which are more easily measured, is potentially important for alloy

selection and screening purposes.

Notched specimens. Previously,® we discussed the fact that at 10® cycles

the notched specimen fatigue strength ratio NFSR (NFS/NTS) is nearly con- stant (0.22 ± 0.01) for these four alloys. This implies that tensile tests

of notched specimens, which are relatively easy to perform, may be used to

predict the fatigue resistance of notched specimens at 10® cycles. More

data are needed to confirm the value of the NFS/NTS ratio as a screening

parameter for notched specimen fatigue, and to ascertain the range of applicability of this parameter for various alloys, temperatures, and stress

levels

.

Figure 7. The relation between FS and UTS for the four aluminum alloys

(4 - 295 K).

1148

380
Alloy Selection

Many aluminum alloys chat remain ductile and tough at low temperatures

might be used in cryogenic structures but each alloy has advantages and

disadvantages that influence material selection for specific applications.

Availability, weldability, and cost must always be considered in material

selection for a practical application; yet, considering the mechanical

properties alone we can conclude that none of the studied alloys is universally superior. For example, at 76 K, 7075-T6 offers the highest FS , YS,

and UTS in the unnotched condition, 7475-T761 offers the highest tensile

strength in the notched condition, and 2219-T87 offers the highest fracture

toughness (judging from the NTS/YS ratio) . At 4 K, alloy 2090—T8E41 offers

the highest notched specimen fatigue strength and relative fatigue strength

for the cycle life of interest.

Alloy 7075—T6 would be the best candidate for unnotched components when

service is limited by cryogenic fatigue. However, alloy 2219-T87 is usually

chosen for cryogenic aerospace applications because it is readily weldable

and also has good fatigue resistance at 76 K and below. If notches are present and stress concentration factors are high, 2090—T8E41 may be

preferred. For applications at 4 K requiring Nf ^ 2.5 x 10 s cycles, 2090-

T8E41 is a good candidate for either smooth or notched components. Alloy

2090 is a relatively new material that offers special advantages: a 7% lower

density, a 10% higher modulus, and lower fatigue notch sensitivity factor

than conventional aluminum alloys. Additional tests of unrecrystallized

2090, or of other Al-Li alloys, are of interest for future studies.

CONCLUSIONS

1. For the unnotched condition, alloy 7075—T6 has the highest fatigue

resistance among the studied alloys between 295 and 4 K. The rank of

the other alloys varies with the test temperature and cycle life of

interest.

2. For unnotched specimens, the FS increases strongly with decreasing

temperature, but the temperature effect diminishes at lower stress

levels. For lives of 10 6 or 10 6 cycles, we find useful linear correlations between FS and the UTS when the data at all three temperatures are pooled.

3. At 76 K, the fatigue resistances of all four alloys are substantially

reduced by the notch; but the notched specimen fatigue strength of 2090-

T8E41 is better than that of the other three alloys when the notched

specimen fatigue values and fatigue notch sensitivity factors (q) are compared at 76 K and a life of 10 s cycles.

4. For notched specimens of all four alloys, the ratio of the NFS

(notched- specimen fatigue strength) at 10 6 cycles to the NTS (notchedspecimen tensile strength) is nearly constant (0.22 ± 0.01), which

suggests that it may be possible to predict the NFS from the NTS for

high strength aluminum alloys, under certain specified conditions.

ACKNOWLEDGMENTS

This work was sponsored by the Superconducting Super Collider Central

Design Group, Universities Research Association, and by the Department of

Energy, Office of Fusion Energy.

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381
REFERENCES

1. J.W. Morris, Jr., and J. Glazer, Mechanical Behavior of Aluminum Alloys

at Cryogenic Temperatures, in: Cryogenic Materials ' 88 , Vol. 2 , Int. Cryo. Maters. Conf . , Boulder CO, R.P. Reed, Z.S. Xing, and E.W.

Col lings , Eds., 1989, pp. 713-726.

2. C. Peters, K. Mirk. A. Wandestforde and C. Taylor presented at the

Tenth International Conference on Magnet Technology. Boston,

September 21—25, 1987.

3. R.L. Tobler and D.T. Read, Fatigue Resistance of a Uniaxial SGlass/Epoxy Composite at Room and Liquid Helium Temperature,

J. Compos. Mater. 10:32—43 (1976).

4. J.W. Kaufman, G.T. Sha, R.F. Kohn, and R.J. Bucci, in: Cracks and

Fracture, ASTM STP 601, Amer. Soc . Test. Maters., Philadelphia,

1976, pp. 169-190.

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

Fracture, and Fatigue Properties of Notched Aluminum Alloy Sheets at

Liquid Nitrogen Temperature, in the proceedings of the fifth interinternational conference, "Aluminum-Lithium Alloys V", Williamsburg,

VA, March 27-30, 1989, to be published.

6. R.L. Tobler, J.K. Han, and R.P. Reed, Fatigue Resistance of a 2090-

T8E41 Aluminum Alloy at Cryogenic Temperatures, in: Cryogenic

Materials '88, Vol. 2, 1988 Int. Cryo. Mater. Conf., Boulder, 1988,

pp. 703-712.

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MBT*114A U.S. DEPARTMENT OF COMMERCE

(REV. 9-m NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY

BIBLIOGRAPHIC DATA SHEET

1. PUBLICATION OR REPORT NUMBER

NISTIR 3944

2. PERFORMER! ORGAMZATtON REPORT NUMBER

X PUBLICATION DATE

September 1990

4. TITLE AMD SUBTITLE

MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES-XIII

S. AUTHOR(S)

R.L. Tobler and R.P. Reed, editors

6. PERFORMING ORGANIZATION (IF JOINT OR OTHER THAN NIST, SEE INSTRUCTIONS)

U.S. DEPARTMENT OF COMMERCE

NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY

Boulder, CO 80303

7. CONTRACT/QRANT NUMBER

L TYPE OF REPORT AND PERIOD COVERED

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (STREET, CITY, STATE, ZIP)

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

10. SUPPLEMENTARY NOTES

Previous reports in this series: NISTIR 89-3931; NBSIR 88-3082; NBSIR 87-3067;

NBSIR 36-3050; NBSIR 85-3025; NBSIR 84-3000; NBSIR 83-1690; NBSIR 82-1667; NBSIR 81-1645;

.NBSIR 80-1627; NBSIR 79-1609; NBSIR 78-084.

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 research to determine the properties of materials

that may be used in cryogenic structures for the superconducting magnets of fusion energy

power plants and prototypes. Its purpose is to facilitate their design and development.

The program was developed jointly by the staffs of the National Institute of Standards

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

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

laboratories through cooperative agreements with NIST.

Research results for 1989-90 are presented in technical papers under four headings that

reflect the main program areas: Structural Alloys, Welding, Test Standards, and Technology

Transfer. Objectives and research highlights are summarized in the introduction to each

program area.

12. KEY WORDS (6 TO 12 ENTRIES; ALPHABETICAL ORDER; CAPITALIZE ONLY PROPER NAMES; AND SEPARATE KEY WORDS BY SEMICOLONS)

austenitic steels; cryogenic properties of materials; low temperature; mechanical

properties; stainless steels; standards; structural materials; velds

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