A11103 Qomia

NBSIR 86-3050

“"111

1 TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES - IX

TEMPERATURE

To

Department of Energy

Office of Fusion Energy

'"^-^Hlngton, D.C. 20545

Prepared by

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

100

.056

86-3050

1986
NATIONAL INSTITUTE OF STANDARDS &

TfiCHNOLOGY

Research InfcHinatioii Center

Gaithersburg, MD 20899
NBSIR 86-3050 UC 20b, c

/y^sv

TECHNICAL REPORTS GlCioc,

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES - IX

Edited by r^.

^

R.P. Reed ^ ^

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

May 1986

Sponsored by

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

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

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

CONTENTS

SUMMARY V

ORGANIZATIONAL CONTACTS vi

PROGRAM DESCRIPTION vii

HIGHLIGHTS OF RESULTS 1 STRUCTURAL ALLOYS

Structural Alloys Program 11 Summary of Mechanical Property Measurements of Structural Alloys . 12

Alloy 31 6LN for Low Temperature Structures:

A Summary of Tensile and Fracture Data 15

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

Strength and Toughness of AISI 30l1 and 316 at 4 K 27

N. J. Simon and R. P. Reed

Near-Threshold Fatigue Crack Growth Behavior of

AISI 316 Stainless Steel 43

R. L. Tobler

The Toughness of Austenitic Stainless Steels at 4 K 53

P. T. Purtscher and R. P. Reed

Absence of Stretch Zones in Austenitic Stainless

Steels Fractured at Cryogenic Temperatures 89

P. T. Purtscher

Low Temperature Deformation of Copper and an Austenitic Stainless Steel 95

R. P. Reed and R. P. Walsh

Low Temperature Tensile Characteristics

of Copper and Aluminum 107

R. P. Reed and R. P. Walsh

Effect of Prior Deformation on the 76-K Fracture Toughness of

AISI 304L and AWS 308L Stainless Steels 1 4l P. T. Purtscher and D. T. Read

Carbon and Nitrogen Effects on the Elastic Constants

of a Stainless Steel at 4 K 159

H. M. Ledbetter, M. W. Austin, and S. A. Kim

Effects of Interstitial Carbon and Nitrogen on

the Low-Temperature Magnetic Properties of FCC

Fe-Cr-Ni Alloys 167

E. R. Jones, Jr., T. Datta, C. Almasan,

and H. M. Ledbetter

ill
STRUCTURAL ALLOYS (continued)

Proposed Standard Method for Tensile Testing of

Structural Alloys at Liquid Helium Temperature 177

Proposed Standard Method for Fracture Toughness Testing

of Structural Alloys at Liquid Helium Temperature 195

Minutes of Low Temperature Structural Materials and Standards

Workshop 209

WELDING

Welding Program 213

Summary of Mechanical Property Measurements of Welds 21^

Manganese and Nitrogen in Stainless Steel SMA Welds

for Cryogenic Service 217

C. N. McCowan, T. A. Siewert, and R. P. Reed, and F. B. Lake

Strength-Toughness Relationship for Austenitic

Stainless Steel Welds at 4 K 239

R. L. Tobler, T. A. Siewert, and H. I. McHenry

Tensile and Fracture Properties of an Fe-1 8Cr-20Ni~5Mn-0. 1 6N

Fully Austenitic Weld Metal at ^ K 2^7

C. N. McCowan, T. A. Siewert, and R. L. Tobler

Tensile, Fatigue, and Fracture Properties of an Fe-l8Cr-l6Ni-6.5Mn-2.4Mo Fully Austenitic SMA Weld at 4 K 257

R. L. Tobler, R. E. Trevisan, H. 1. McHenry,

P. T. Purtscher, and T. A. Siewert

NONMETALLICS

Nonmetallics Program 279

Irradiation Effects on Organic Insulators 28l

M. B. Kasen

Screening the Performance of Organic Insulators

under Cryogenic Neutron Irradiation 293

M. B. Kasen and R. B. Stoddard

Young Modulus and Internal Friction of a Fiber-Reinforced Composite 301

H. M. Ledbetter, M. Lei, and M. W. Austin

Edge Stresses in Woven Laminates at Low Temperatures 307

R. D. Kriz

TECHNOLOGY TRANSFER

Technology Transfer Program 327

Cryogenic Structural Materials Workshop 329

Information Retrieval Services 337

iv
SUMMARY

This report contains results of a research program to

produce material property data that will facilitate design and

development of cryogenic structures for the superconducting

magnets of magnetic fusion energy power plants and prototypes.

The program was developed jointly by the staffs of the National

Bureau of Standards and the Office of Fusion Energy of the

Department of Energy; it is managed by NBS and sponsored by DOE.

Research is conducted at NBS and at various other laboratories

through subcontracts with NBS.

Research results for 1985 are summarized in an initial

"Highlights of Results" section and reported in detail in the

technical papers that form the main body of this report. The

technical papers are presented under four headings reflecting

the main program areas: Welding, Nonmetallics ,• Structural

Alloys, and Technology Transfer. Objectives, approaches, and

achievements are summarized in an introduction to each program

area.

Note : Certain commercial equipment, instr'uments , or materials

are identified in this report to specify the experimental

procedure adequately. In a few instances, company names are used to identify the source of specific research. In no case

does such identification imply recommendation or endorsement by

the National Bureau of Standards, nor does it imply that the

material or equipment identified is necessarily the best

available for the purpose.

V
ORGAHUAT/OHAL COHTACTS

The staff contributing to the various major aspects of the

program during fiscal year 1986 are listed below so that specific

technical questions may be directed to them.

Department of Energy, Office of Fusion Energy, Washington, D . 0 . 205 ^!

Program Monitor V. K. Der (301 ) 353-5736

(FTS) 233-5736

National Bureau of Standards, Boulder, Colorado 80303

Program Manager R. P. Reed (303) 497-3870

(FTS) 320-3870

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

S. A. Kim (303) 497-3504

H. M. Ledbetter (303) 497-3443

P. T. Purtscher (303) 497-5789

R. P. Reed (303) 497-3870

N. J. Simon ( 303 ) 497-3687

R. L. Tobler ( 303 ) 497-3421

R. P. Walsh (303) 497-3537

Welding C. N. McGowan (303) 497-3523

H. I. McHenry (303) 497-3268

P. T. Purtscher ( 303 ) 497-5789

R. P. Reed (303) 497-3870

T. A. Siewert (303) 497-3523

R. L. Tobler (303) 497-3421

D. P. Vigliotti ( 303 ) 497-3222

Nonmetallics M. W. Austin (303) 497-3504

M. B. Kasen ( 303 ) 497-3558

R . D . Kr i

z

(303) 497-3547

H. M. Ledbetter (303) 497-3443

M. Lei (303) 497-3504

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

N. J. Simon (303) 497-3687

vi
PROGR/IM VESCRIPTIOH

The overall objective of the program is to assist in the design, con- struction, and safe operation of low-temperature magnetic fusion energy (MFE)

systems, especially superconducting magnets, through effective materials

research and materials technology transfer. The specific steps taken to

achieve this objective are: (1) evaluation of low-temperature materials

research needs specific to MFE devices; (2) development and monitoring of a research program to acquire the necessary data; and (3) 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 recent 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 austenitic-steel elastic

constants also contributed to this objective. Through interactions with

low-temperature design, construction, and measurement programs, such as the

Large Coil Project and MFTF-B, we are aware of new problems as they arise.

Research projects contributing to the second objective are outlined in the

following table. The results appear later in this report. Again, emphasis

has been on research oriented toward the establishment of test standards.

The third objective is satisfied, in part, by these annual reports and by the

series of workshops on Materials at Low Temperatures. Since 1982, handbook

pages presenting the available data for specific materials have been distributed to those members of the magnetic fusion energy community involved with

low-temperature materials.

vii
NBS-DOE RESEARCH PROGRAM - MATERIALS AT LOW TEMPERATURES

R. P. Reed, Program Manager

Structural Alloys Welding

R. P. Reed, Leader T. A. Siewert, Leader

• Development of strong, tough

austenitic steels

• Characterization of high-strength,

high-conductivity copper and

copper alloys

• Development of low-temperature test

standards

• Development of improved filler metals

for low-temperature service

• Detection and evaluation of defects

in stainless-steel weldments

• Evaluation of mechanical properties

of weldments at low temperatures

Nonmetallics

M. B. Kasen, Leader

Technology Transfer

N. J. Simon, Leader

• Measurement, modeling, and understanding properties of composites

• Development of specifications for

insulating laminates

• Development of test methods

for materials selection

• Evaluation of data on austenitic

steels, copper, and copper alloys

• Organization of cryogenic structural materials workshop

• Dissemination of materials properties data to DOE and the fusion

energy community

viii
HIGHLIGHTS Of RESULTS

1

HIGHLIGHTS Of RCSUl TS

STKUCTUKAl ALLOYS

Strong, Tough Alloys

The tensile-deformation characteristics and effect of strain rate were studied on relatively pure CDA 101 copper, 1100 aluminum, and solid-solutionstrengthened AlSl 310 . Tensile strain rate was varied between 2 x 10 ^ and

2 X lO"^ s””' at temperatures ranging from ^ to 295 K, Tensile stress-strain

and strain-hardening curves were determined for these temperatures. The

effect of strain-rate changes on tensile flow strength was measured from

strains near 0.002 (yield strength) to over 0.300. The data reflect three

distinct ranges of face-centered cubic, polycrystalline plastic deformation

whose characteristics depend on the solute content.

Provisional equations for the yield strength of AISI 304- and 3l6~type

alloys at 4 K as a function of nitrogen content and grain size were developed. Provisional equations for the 4-K fracture toughness, Kj^(J), of

these alloys are presented in which the effects of yield strength, nitrogen

and manganese content, and inclusion spacing are expressed quantitatively.

The equations result from regression analyses of a matrix of NBS measurements

that includes more extensive numerical data on alloying, refining, and processing parameters than have previously been available.

For structural use in superconducting magnet applications, alloy AISI

316 LN offers higher strength than AISI 304LN with equivalent toughness.

Tensile and fracture toughness data at 4 K were measured for a variety of

AISI 316 - and 3l64,N-alloy specimens obtained from suppliers in Japan, Europe,

and the United States.

Five elastic constants of nine AISI 304-type stainless steel alloys were studied between 295 and 4 K. The carbon-plus-nitrogen contents of these

alloys ranged from O.O 8 O to 0.359 wt.^. Longitudinal, Young’s, shear, and

bulk moduli and Poisson's ratio were determined by a pulse echo ultrasonic

method. Except for the bulk modulus, all the elastic stiffnesses increased

slightly with increasing interstitial content.

The role of alloying and inclusions in changing the toughness of

austenitic stainless steels was assessed. The toughness was evaluated

according to ASTM E 81 3 standard for elastic-plastic fracture. The alloys

were chosen carefully so that the effect of substitutional alloying elements

(nickel, manganese, molybdenum) could be studied; other variables that

influence the strength and toughness were kept to a minimum. This information is valuable in designing alloys that will meet specifications for

service at 4 K with alloying elements that are cost effective.

3
strong, Fatigue-Resistant Alloys

A new facility is in use that facilitates the difficult and timeconsuming measurement of threshold fatigue. It is capable of measuring

fatigue-crack growth rates as low as one angstrom (less than one lattice

spacing) per load cycle in the temperature range 295 to 4 K. Near-threshold

fatigue-crack growth rates and threshold stress Intensity factors are important because they are used to determine the stress limits below which

existing cracks cease to grow.

The near-threshold fatigue behavior of an AISI 316 alloy was characterized using this new, fully automatic fatigue test apparatus. Significant

differences in the near-threshold behavior at temperatures of 295 and 4 K were observed. At 295 K, where the operationally defined threshold at

10”^^ m/cycle is insensitive to stress ratio and strongly affected by crack

closure, the effective threshold stress Intensity factor is about 4.65

MPa*m”''^^ at R = 0.3 and 6.1 MPa*m^'^^ at R = 0.1.

Codes and Standards

A proposed standard method for tensile testing of structural alloys at

4 K was developed. Procedures used in displacement-controlled tensile

testing of structural alloys and weldments at this temperature are described.

The format corresponds closely to ASTM standard designation E 8-83 entitled

"Standard Methods of Tensile Testing of Metallic Materials." The contents

include modifications to accommodate the special needs of cryogenic testing.

A proposed standard method for fracture toughness testing of structural

alloys at 4 K was also developed. It covers the equipment, procedures, and

data analysis used for determining the fracture toughness of structural

alloys and weldments at this temperature. The procedures incorporated

represent adaptations of the and Jj^ plane strain fracture toughness

tests, as described in ASTM Standards E 399-83 and E 8l 3~8l , respectively, to

accommodate the specific requirements of cryogenic testing.

4
WELDING

A series of welds based on an AISI 308-type composition was used to

clarify the effects of manganese (1.5 to 10 wt./5), nitrogen (0.04 to 0.26

and molybdenum (0 to 4 wt.^) on SMA weld properties for cryogenic

service. The 4-K yield strength was found to be strongly increased by

nitrogen, but only weakly increased by manganese. The ?6-K lateral expansion

had a more linear correlation to the strength and alloy content than the CVN

absorbed energy. Present equations for predicting ferrite numbers were found

to be quite inaccurate for welds with high manganese and nitrogen levels.

Improved predictive equations were determined.

A fully austenitic welding electrode was evaluated at 4 K. A 1.2-mmdiameter electrode with a composition of Fe-1 8Cr-20Ni-5Mn-0 . 1 6N was used to

produce a gas-metal-arc weld in 25~mm-thick AISI 304-type stainless steel

plate. The yield strength (1015 MPa) and fracture toughness (203 MPa*m"'

compare favorably with existing base metal properties for AISI 304-type

alloys.

A fully austenitic butt weld, produced by shlelded-metal-arc welding

with the composition Fe-l8Cr-l6Ni-6.5Mn-2.4Mo, was evaluated at 4 K. An

average fracture toughness, Kj^(J), of 191 MPa*m^^^ was measured on the

25~mm-thick specimens. A yield strength 842 MPa was measured, and the

tearing modulus was determined to be 50. The fracture toughness was better

than common AWI 308- and 3'16-type welds, but poorer than base materials of

comparable strength.

The strength-toughness relationships for welds evaluated at NBS over the

past decade were collected and compared with the corresponding base material

properties. An inverse linear correlation was found between yield strength

and fracture toughness for AWI 308- and 31 6-type welds at 4 K. The AWI 308-

and 316 -type welds had a significantly lower toughness than base materials of

comparable strength. Newer, fully austenitic alloys produced welds with

better toughness.

5
NOHMETALUCS

standard Specimen Development

Recognizing the unique demands placed on specimens to be used in cryogenic radiation research, we continued development of the technology for

producing specimens combining the attributes of low cost, a high degree of

characterization, small size, flexibility in component selection and treatment, and adaptability to many test procedures. Continuing our emphasis on

the basic 3 . 2-mm-diameter rod configuration, we have extended our production

capability to include a solid component epoxy resin system (G-llCR) and a bismaleimide polyimide. Materials required to perform an initial irradiation

test run in the NLTNIF reactor have been produced and sent to ORNL. These

consist of three epoxy resin systems and one bismaleimide system, each

produced as unreinforced resin and as composites reinforced with 48 volume

percent glass fiber.

Standard Specimen Test Methods

We developed and validated the apparatus required to assess the influ- ence of cryogenic neutron irradiation on mechanical property degradation of

highly characterized test specimens. Emphasis was placed on test methods

sensitive to degradation of fiber-matrix interfacial properties. We have

successfully modified the conventional short-beam shear test procedure to

accept the standard rod specimens. We have developed the complete capability

for efficiently conducting torsional shear tests at cryogenic temperatures

and have demonstrated the sensitivity of this test method to assessing interfacial property changes. This test method has not been applied previously to

composite materials at cryogenic temperatures. We have also developed a simple, novel method of obtaining quantitative information on the longitudinal fracture strength, of the standard composite rod specimens at

cryogenic temperatures. This is also a new development in cryogenic composite testing technology.

International Cooperation

The cooperative international program directed toward development of

radiation-resistant, organic-matr ix insulators is continuing. The group at

ISIR, Osaka University, Japan, headed by Prof. T. Okada, continues to produce

and test standard composite specimens of configuration identical to those

developed at NBS. Prof. Okada will provide an additional series of specimens

to be irradiated at the NLTNIF facility at ORNL during 1986 and will provide

personnel from his organization to conduct the subsequent test work at

NBS-Boulder

.

6
Property Determination and Modeling

We developed a finite element model that predicts stress singularities

near the stress-free edges of laminates. The edge stresses parallel to the

surface are nonzero and large enough to cause the formation of ply cracks and

delaminations. In the laboratory, we rarely observed delamination initiation

at the edges of woven laminates at low temperatures. Finite element theory

explains the influence of weave geometry on these edge stresses and the

amount of stress required to initiate edge de lamination.

By a kilohertz-frequency resonance method, we determined the Young's

modulus and internal friction of a uniaxially fiber-reinforced composite.

The composite comprised glass fibers in an epoxy-resin matrix. We studied

three fiber contents: 0, 41, and 49 vol.^. Young's modulus fit a linear

rule of mixture. The interhal friction fit a classical free-damped-oscillator

model where one assumes a linear rule of mixture for three quantities: mass,

force constant, and mechanical-resistance constant.

7
TBCHHOLOey TRAHSfER

The workshop on Materials at Low Temperatures that has been organized by

NBS in past years took place at LBL, in Berkeley, California, October 2-4,

1985 , under the sponsorship of LBL and LLNL. The workshop brought together

material specialists and magnet designers to report recent research and to

identify and discuss structural materials problems in superconducting magnet

design and construction for fusion energy systems. The proceedings of this

informal workshop were not published; however, a copy of the program and an informal summary of the proceedings may be found in the Technology Transfer

section of this report.

Handbook and supporting documentation pages on mechanical properties of

AISI 304 stainless steels with nitrogen content ^0.1 wt.5^ were distributed

in preliminary draft format in March 1985. Properties covered included:

tensile strength, yield strength, elongation, reduction of area, stressstrain curves, elastic constants, impact energy, fracture toughness, stressand strain-controlled fatigue life, and fatigue crack growth rate over the

temperature range 3 OO to 4 K. The scope of the handbook has been expanded to

include analysis and correlation of AISI 304 properties with respect to

interstitial carbon and nitrogen content. Since very little physical

property data on high nitrogen 304 alloys exist, an inclusive survey of

physical property data from 3 OO to 4 K on all 304 stainless steel alloys was undertaken. A preliminary draft of handbook and supporting documentation

pages on the physical properties of AISI 304 alloys was distributed in late

1985 . Properties covered were: specific heat, thermal conductivity, thermal

expansion, electrical resistivity, magnetization, and magnetic susceptibility. Preliminary work on handbook pages covering properties of CDA 102

copper and selected copper alloys is in progress.

8
STRUCTURAL ALLOYS
.'»v,
STRUCTURAL AUOyS PROGRAM

LEADER: R. P. Reed

CONTRIBUTORS: M. W. Austin, S. A. Kim, R. D. Kriz, H. M. Ledbetter,

P. T. Purtscher, Y. Shindo, N. J. Simon, R. L. Tobler,

R. P. Wa Ish

OBJECTIVES

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

magnets at 4 K

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

pulsed superconducting magnets at 4 K

• Development of effective codes and standards related to low-temperature

property measurements and structural design

RESEARCH PAST YEAR (1985)

• Low-temperature strain-hardening and strain-rate effects of copper,

aluminum, and AISI 310

• Threshold fatigue crack growth behavior of austenitic stainless steels

• Drafts of proposed standard methods for tensile and fracture toughness

testing at 4 K

• Summary of strength and toughness characteristics of AISI 304- and

31 6-type alloys

• Elastic properties of austenitic stainless steels at low temperatures

RESEARCH THIS YEAR (1986)

• Low-temperature strain-hardening and strain-rate effects of austenitic

stainless steels

• Cooperative program with industry on high-nitrogen, high-molybdenum

austenitic stainless steels

• Continued development of standard methods for tensile and fracture

toughness testing at 4 K by R. L. Tobler at Tohoku University, Sendai,

Japan

• Cooperative program with LLNL and MIT on high-strength, high-conductivity

copper alloy properties

• Creep studies of austenitic stainless steels and copper at low

temperatures

• Study of plastic fracture in austenitic stainless steels at 4 K

• Study of the effects of magnetic field on the crack tip singularities in

ferromagnetic materials; participation of Y. Shindo, Tohoku University, in

this study

11
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS

Alloy

(Designation)

Country

of

Supplier

Tensile

295-44 K Elastic295-44 K

Fracture Toughness

295, 76, 4j K

Fatigue Crack

Growth Rate

295, 76, 4; K

Fatigue

S-N

44 K Report Volume*

(page)

AUSTENITIC STEELS

Fe-5Cr-25Mn-0.2C Japan VII VIK65)

Fe-5Cr-26Mn Japan VI VK181

)

Fe-1 3Cr-19Mn USSR III VI 111(79), VI(441 )

Fe-13Cr-22Mn-0.2N Japan VII VIK65)

Fe-1 6Cr-8Ni-8Mn U.S. VIII VIII VIIK151 ,181

)

Fe-1 7Cr-l8Mn-0.5N U.S. VII, VIII VIII VIK65) ,VIII(181 )

Fe-17Cr-9Ni-8Mn U.S. III 111(91

)

Fe-17Cr-13Ni-2Mo

(316) U.S. I, II, III,

IV, VI, VIII

VI I, II, VIII

II, VIII I 1(1 5, 71 ), 11 (79), 111 (449, 105,1 17), IV( 1 447 ),V(185),

VI(87),VI(157),VIII(181 , 209,251

)

Fe-1 7Cr-1 3Ni-2Mo

(316LN)

U.S. II II II 11(79)

Fe-l8Cr-3Ni-l 3Mn U.S. I, VI, VIII III I, VIII I, VI, VII 1(93), 111(91 ),VI(53),

VII(85),VIII(167,219)

Fe-19Cr-(6-15)Ni44Mn-(0.1-0.3)N USSR VIII VIII VIIK123)

Fe-19Cr-9Ni

(3044 ) U.S. I, II, IIIIV, VI, VII

I, II, III

IV, VI, VII

I, II, III,

IV

II, III, IV I 1(15,71 ,21 3), 11(79, 1 449, 175) ,111(15,91 ,105,

117),IV(37,101 ,203,215,

227),V(71),VI(73),VII

(13), VII( 447 ),VII(75),

VIK157)

Fe-1 9Ci—9Ni

(3044L) U.S. II, III,IV, VIII

I, II, IV III, IV VIII

III, IV VIII

1(213), 11(79, 123, 175),

IV(37,101 ,131 ,215,302),

VIIK 181 ,229,251 )

Fe-19Cr-9Ni-(1-8)Mn

(3044LN) U.S. IV, V, VIII V IV, V, VIII IV(77),V(15,59,189),

VIIK181)

Fe-19Cr-9Ni-N

(3044N) U.S. I, III,

IV IV I, III, IV I, III, IV I(93),III(15),IV(37,101 ,203,215)

Fe-19Cr-9Ni-N

(304LN)

U.S. II, III, IV,V,VI,

VIII

IV, VIII II, III, IV, VIIIII, III, IV, VIIIII 11(35, 79), 111(15), IV (37,101 ,203),V(29),VI

(113),VIII(181 ,207,229)

Fe-19Cr-10Ni-

(0.03-0.29)N-

(0.03-0.09)0

U.S. III, VIII VIII III, VIII III 111(15) ,VIII( 1 445, 1 8 1 )

Fe-20Cr-7Ni-8Mn

(216) U.S. VIII VIII VIIK151 ,181 )

MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-88A (I); NBSIR 79-1609 (II): NBSIR 80-1627 (III); NBSIR 81-16145 (IV); NBSIR 82-1667 (V); NBSIR 83-l690(VI);

NBSIR 844-3000 (VII); NBSIR 85-3025 (VIII).

12
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

A1 loy

( 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

S-N

4 K Report Volume*(page)

Fe-20Cr-8Ni-2Mn

(304HN)

U.S. VIII VIII VIIK 151 ,181 )

Fe-20Cr-1 6Nl-6Mn USSR V,VI V V(29,213),VI(113)

Fe-21Cr-6Ni-9Mn U.S. I, VIII III, VI I, VIII I 1(15,71 ),III(91 ), VI(157),VIII(151 ,181

)

Fe-21Cr-12Ni-5Mn U.S. I, II III I. II I, II 1(93), 11(79), 111(91 )

Fe-25Cr-21Ni

(310)

U.S. V, VI, VII, VIIIIII, VI V.VIII 111(105,1 17), V(71 ,145),

VI(73),VII(65),VIII(181 )

Fe-19Ni-9Co

( 200-300 grades)

U.S. IV IV(237)

Reviews of Alloys IV,V,VI IV,V,VI IV,V,VI IV,V,VI

,

VII

IV, VI IV(17,257),V(171 ), VK11 ),VI(127),VI(157),

VIKIO3 )

ALUMINUM ALLOYS

Al-6Cu-0.2Mn-0.1Fe

(5083-0)

Al-4Mg

U.S.

U.S. IV

II II II II

II

11(19)

11(35), IV(185)

(3083-0)

SUPERCONDUCTORS

Nb-i^STi U.S. Ill 111(133)

•MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-88i1 (I); NBSIR 79-1609 (II): NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-l690(VI);

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII).

13

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15 . -s

ALLOY 316LN FOR LOW TEMPERATURE STRUCTURES:

A SUMMARY OF TENSILE AND FRACTURE DATA*

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

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

For structural use in superconducting magnet applications, alloy 316LN

offers higher strength than 304LN with equivalent toughness. Nitrogen additions permit one to obtain a range of yield strength levels with little added

cost. Tensile and fracture toughness data at 4 K are presented for a variety

of 316 and 316 LN alloys. The specimens were obtained from suppliers in Japan,

Europe, and the United States and were tested at NBS. Elastic-plastic fracture toughness was determined from J-integral measurements.

INTRODUCTION

A number of alloys suitable for structural use at low temperatures have

been studied at the National Bureau of Standards. These austenitic steels,

containing nitrogen additions to increase their tensile yield strength at 4 K,

include: 304-type (Fe, l8-20Cr, 8-IONi, 1-2Mn, 0.10-0.20N), high-manganese

(Fe, 13-21Cr, 1-7Ni, 9-25Mn, 0.20-0.30N), and 316-typb alloys (Fe, 17~19Cr,

11-14N1, 1-2Mn, 2

-

3M0 , 0.10-0.25N). These alloys generally have sufficient

strength, toughness, and weldability required of the principal structural

material for high-field superconducting magnets. Tensile and fracture toughness data of the 31 6-type alloys are reported and reviewed i,n this paper.

PROCEDURES

The chemical composition, hardness, and type and number of inclusions for

each 316 and 3 I 6 LN alloy that we have studied at 4 K are listed in Table 1.

Only alloys B and D are from laboratory heats; the remainder are from commercial heats that have been used in superconducting magnet construction. The

alloys were prepared in Europe, Japan, and the United States. The inclusions

were analyzed using a scanning electron microscope with energy-dispersive

X-ray analyses. The sensitivity of particle detection was about 0.5 pm. The

typical size of inclusions was 1 .0 pm. The inclusions were analyzed on polished surfaces, not on the fracture surface. Stringers of manganese sulfide

were observed in all alloys except B and D; the content of other inclusions

Included Al, Ti, Or, Mg, Ca, and Si.

^Submitted to Proceedings of the 11th International Cryogenic Engineering

Conference.

15
All tensile specimens were about 6 mm in diameter with a reduced section

about ^0-mm long. Strain-gage, clamp-on extensometers with a 25-mm gage

length were used to record strain. All tensile procedures and experimental

sensitivities and accuracies have been previously discussed.^

To determine the fracture toughness at ^ K, J-integral measurements were employed. Our technique uses compact tensile specimens and follows ASTM E 813

procedures.^ To convert the measured

'^Ic to Kjq(J), the expression

J Ic

[Kj^(J)]2

E

( 1 )

was used. Here is the measured critical value of the J-integral and E is

Young’s modulus (207 GPa at 4 K) . Our previously reported values^ of Kj^(J)

for alloys E, F, and G were calculated using the Poisson's ratio factor

(1 - v^) in the denominator of Eq. (1). These values are adjusted in this

paper to Eq. (1) formulation. In all reported measurements, we used 25-mmthick compact tensile specimens with a strain-gage, clip-on extensometer to

measure crack-opening displacement.

RESULTS AND DISCUSSION

(a) Strength

The tensile yield strength (Oy, 0.2^ offset) at 4 K of the completely

austenitic 316-type alloys depends primarily on two variables: grain size and

N content. The grain sizes and yield strengths of the alloys tested are given

in Table 2; their N content, in Table 1.

Grain size was related to the yield strength for alloy G in a previous

study by Tobler et al.'^ The dependence of yield strength on grain size

(obtained by recrystallization and annealing for different times at 1323 K)

of alloys B and D is shown in Figure 1. Here d (d is the average grain

diameter in pm) is linearly related to yield strength, as predicted from the

well-known Hall-Petch relationship

Oy = Oq + md~^^^, (2)

with Oq the single-crystal yield strength and m a proportionality constant.

The strong dependence of yield strength on N content is shown in Figure 2. Here, the yield strengths are normalized to a grain size of 50 ym,

using the average dependence of alloys B and D (Fig. 1). As an interstitial,

N contributes substantially more to the low temperature strength than other

alloying elements.

Our linear least -squares computer program was used to analyze the entire

data set of Table 2 for the effect of grain size and N content in wt./5 ([N]) on Oy at 4 K. The equation developed.

16
Oy (MPa) = 381 + 2778 [N] + 81 1 d (3)

has a standard deviation (SD) of 55 MPa for Oy, and SD’s of 3^, 193, and 199

for the constant term and the coefficients of [N] and

The dependence on [N] is not as large as that reported by Reed and Simon^

for 30^1 and 3Qi1LN alloys:

Oy (MPa) = 325 + 3^00 [N]. (4)

In their analyses, C content was included as a variable, the dependence on grain size was not considered, and data from other laboratories were also

included. The higher value of the constant term in Eq. (3) was expected,

since added strengthening is obtained from Mo addition in the 316 alloys.

In both studies, the concentrations of the other elements were held rather

constant: 18-19 wt.5^ Cr , 8-12 vjt.% Ni, 1-2 wt.^ Mn, 2-2.5 wt./S Mo, and

0.03 wt.$ C. Thus, by maintaining this composition one can use Eq. (3) to

predict yield strength at 4 K on the basis of grain size and N content.

The addition or control of N content is less expensive than control of

grain size of strong, thick-section plate, which demands large-capacity rolling mills, accurate temperature control, and perhaps, facilities for rapid

cooling from the hot-rolled condition to near room temperature. However, the

solubility of N in a 316-type alloy may be restricted to slightly over

0.20 wt.?. Solid-solution strength increases beyond the contribution of 0.20

vit.% N must come from more expensive Mo additions.

(b) Toughness

The fracture toughness of austenitic steels at 4 K is primarily dependent on alloy content, strength, and inclusions. Previously, we have shown

that toughness depends on Ni content when inclusions and strength are held

constant, and that toughness depends on N content only through its strengthening effect. ° In this paper we consider the dependence of toughness on strength and inclusion spacing for the 31 6-type alloys.

Most metastable and stable austenitic steels fracture in a ductile manner. Their fracture surfaces exhibit a dimpled appearance, typical of ductile

fracture, in which microvoids form at initiation sites (usually inclusions)

and eventually coalesce into voids. Linkage of the voids produces fracture.

Alloys B - H exhibited ductile fracture in J-integral tests at 4 K. Alloy A,

aged at 873 K from the hot-worked condition, showed irregular dimpled areas combined with planar, brittle-appearing areas. The Kj^(J) data of alloy A are

not used in the following correlations.

The fracture toughness Kj^(J) at 4 K is given in Table 2, and its dependence on tensile yield strength (also at 4 K) is shown in Figure 3 and in the

following equation:

Kj^(J) (MPa-m''/2) = 516 - O. 3 O 6 Oy. (5)

17
The SD values are 70 MPa*m^^2 for Kj^(J), 111 for the constant term, and 0.13

for the coefficient of o . There is an inverse linear dependence of Kjq(J)

on Oy, like that of the 304-alloy series,"^ but there is a sizable SD. However, this variability is from alloy to alloy; when parameters such as the

grain size are varied for one alloy (see Table 2, alloys B and D), the scatter is considerably reduced. This leads one to suspect that inclusions, the

initiation sites for the ductile fracture mechanisms, are responsible for the

large alloy-to-alloy data scatter.

A tabulation of the reciprocal square root of the average number of

inclusions per mm^ for each alloy is included in Table 1. The

types of inclusions vary and are dependent on refining, deoxidizing, and

alloying practices. In Figure 4 we plot the product of the fracture tough- ness and the yield strength versus corresponding to particle density.

The product Kj^(J)*ay is used to normalize the strength effect on toughness,

that is, the strong, inverse dependence of on Oy is neutralized. Previously we have used this quantity, which we have labeled the Quality Index, to

portray the effect of alloying on the relationship between Kj^(j) and Oy.

Vie found a strong dependence of the Quality Index on Ni content. Therefore,

because the Ni content varied from 10.9 to 13-9 wt.% (Table 1), we adjusted

the Quality Index to 12 wt.^ Ni for each alloy using an equation derived from

the data presented by Reed and Purtscher.® There is an approximately linear

dependence of adjusted K^^(J)* 0 y versus (Figure 4) that can be

expressed as

Kj-c(J)- 0 y

(1o3 MPa^-m’’^^) = 74 + 1800 . ( 6 )

The SD values are 28 for Kj^(j)*Oy, 22 for the constant term, and 278 for the

coefficient of

It is reasonable to assume that Inclusion density principally affects

the toughness, not the strength, of our austenitic alloys. Sakamota et al.

have reported data that support this premise. ° Therefore, in Figure 5, we present the same data, with the Kj^(J) data adjusted to n~^^^ = 0.047 and to

12 wt.% Ni with Oy held constant. Normalization of the toughness data of

Figure 3 to constant particle density and Ni content considerably reduces

the data scatter (SD is reduced from 70 to 29)>and the slope more nearly

reflects the dependence of Kjj^(J) on strength

K!^^(J) (MPa-m^^^) = 433 - 0.28 Oy (7)

for an average particle spacing (n ) of 0.047. The SD values are 29 for

Kjc(J), 45 for the constant term, and 0.05 for the coefficient of Oy.

Finally, since we have demonstrated that the toughness varies with number of inclusions and yield strength and have previously shown toughness to

be dependent on Ni content, we performed a linear regression least-squares

fit on these variables for our original 316-type alloy data set (Table 2). We

found that

Kj^(J) (MPa-m''/2) = -,30 - 0.338 Oy + 2250 + 20.2 [Ni]. ( 8 )

18
The SD values are 31 MPaTn^'^^ for 89 for the constant, 0.059 for

the coefficient of Oy, 310 for the n and 5.2 for the [Ni] coefficient.

All independent variables contributed to a reduction of the SD for Kj^(J), and

as indicated, the SD values for the coefficients of the independent variables

were reasonable. This equation is discussed further by Simon and Reed.^

Thus, the requirements for optimum toughness are a minimum inclusion

density (attained by careful refining treatment), optimum Ni content, and

minimal strength. The first two are achieved only by added cost, the third at

the expense of design efficiency. Considering that has a linear in- verse relationship with Oy, the designer has the flexibility of adjusting the

strength-to-toughness ratio through nitrogen content. Higher N increases

strength at the expense of reduced toughness; lower N decreases strength but

increases toughness.

SUMMARY

Two parameters are most effective in controlling the 4-K strength of

316 -type alloys: grain size and N content. Of these, the variation of N

content is the least expensive. Grain-size control of thick plates requires

exceptional rolling capacity.

Toughness is shown to depend on Ni content, inclusion density, and

strength. The addition of Ni is rather expensive, and reduction of inclusions demands careful refining practices. Toughness is sacrificed by increasing strength.

The 316 -type alloys provide excellent flexibility for alloy design of a structural material for use at 4 K. Both strength and toughness may be adjusted by alloy addition, refining practice, and thermomechanical processing.

ACKNOWLEDGMENTS

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

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

R. P. Walsh for conducting all the tensile tests and many of the fracture

toughness tests. L. L. Scull also greatly assisted us by determining interstitial contents in most of the alloys. AMAX Materials Research Center

(P. J. Grobner) and Nippon Kokan Steel, Technical Research Center (C. Ouchi)

supplied a series of laboratory heats. Other alloys were provided by

Lawrence Livermore National Laboratory (E. N. C. Dalder), Oak Ridge National

Laboratory, and Kernforschungszentrum Karlsruhe, Institut fur Technische

Physik (A. Nyilas).

REFERENCES

1. R. P. Reed and R. P. Walsh, Low Temperature Deformation of Copper and an Austenitic Stainless Steel. Advances in Cryogenic Engineering

Materials , vol. 32 , Plenum Press, New York (1986), pp. 303-312; in press.

19
2. Standard Test Method for Jj^, A Measure of Fracture Toughness, 1982 Annual Book of ASTM Standards , Part 10, ASTM E 81 3~8l , American Society for

Testing and Materials, Philadelphia ( 1982), pp. 822-^840.

3. D. T. Read and R. P. Reed, Fracture and Strength Properties of Selected

Austenitic Stainless Steels at Cryogenic Temperatures, Cryogenics

,

vol.

21 (1981), pp. 415-417.

4. R. L. Tobler, D. H. Beekman, and R. P. Reed, Factors Influencing the

Low-Temperature Dependence of Yielding in AISI 316 Stainless Steels,

Austenitic Steels at Low Temperatures , Plenum Press, New York (1983),

pp. 135-157.

5. R. P. Reed and N. J. Simon, Low Temperature Strengthening of Austenitic

Stainless Steels with Nitrogen and Carbon, Advances in Cryogenic Engi - neering—Materials

,

vol. 30, Plenum Press, New York (1984), pp. I 27- 136 .

6. R. P. Reed and P. T. Purtscher, Nickel and Nitrogen Alloying Effects on

the Strength and Toughness of Austenitic Stainless Steels, Advances in

Cryogenic Engineering—Materials , vol. 32, Plenum Press, New York (1986),

pp. 43- 5 O; in press.

7 . R. L. Tobler, D. T. Read, and R. P. Reed, Strength/Toughness Relationship

for Interstitially Strengthened AISI 304 Stainless Steels at 4 K Temperature, ASTM STP 743 , American Society for Testing and Materials, Philadelphia (1982), pp. 250-268.

8. T. Sakamoto, Y. Nakagama, I. Yamauchl, and T. Zaizen, Nitrogen-Containing

25Cr-^13Ni Stainless Steel as a Cryogenic Structural Material, Advances in

Cryogenic Engineering—Materials , vol. 30, Plenum Presd, New York (1984),

pp. 137-144.

9. N. J. Simon and R. P. Reed, Strength and Toughness of AISI 304 and 316 at

4 K, Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures - IX, R. P. Reed, ed.. Internal Report, National Bureau of

Standards, Boulder, Colorado (1986), pp. 27-42.

20
Table 1. Characterization of 316-type Alloys

Alloy

Cr Ni Mn

Chemical Composition (wt.

Mo C N P

%)

S Si

Hardness

Rb

Inclusions*

n

- 1/2

A 16.7 13.7 1 .26 2.7 0.032 0.16 0.01 0.01 0.41 89 U .13

B 18.5 1 1 .14 1 .40 2.5 0.008 0.17 0.002 0.001 0.40 89 0.11

C 16.8 13.9 1 .84 2.2 0.028 0.10 0.022 0.009 0.33 88 0.095

D 18.6 10.9 1 .01 2.2 0.020 0.12 0.002 0.003 0.29 96-102 0.082

E 17.^ 13.9 1 .58 2.5 0.016 0.16 0.021 0.021 0.48 83 0.057

F 17.2 13.5 1 .86 2.3 0.057 0.03 0.024 0.019 0.58 79 0.054

G 17.3 12.2 1 .55 2.2 0.051 0.03 0.026 0.014 0.68 86 0.092

H 18.0 10.5 1.10 2.1 0.01 6 0.14 0.020 0.015 0.54 86 0.048

*n = number of inclusions within an area of 1 mm^

21
Table 2. Mechanical Properties of 3'16-type Alloys at 4 K

Alloy Tensile

(MPa) (MPa)

Elongation

{%)

Reduction

of Area

{%)

Fracture Toughness

(MPa*m ' ^)

Grain Size Condition

(pm)

A 935 1480 36 58 140 63 Aged-837 K

B 1080 -- — -- 220 13 Annealed

942 1590 42 56 — 70 Annealed

924 1560 49 61 320 91 Annealed

917 1570 45 57 310 129 Annealed

C 830 1360 45 53 360 31 Annealed

D 1060 1730 36 49 150 — Annealed

940 1690 41 54 180 1

1

Annealed

880 1620 41 58 190 20 Annealed

870 1640 43 58 — 30 Annealed

840 1630 43 52 — 50 Annealed

1 1 40 1750 34 42 180 44 Hot rolled

1080 1750 38 48 170 51 Hot rolled

1020 1710 37 54 200 53 Hot rolled

1020 1700 38 55 180 41 Hot rolled

E 1010 1390 35 50 210 105 Annealed

F 710 1300 48 57 250 65 Annealed

G 550 1330 52 — 400 — Annealed

500-690 — — — — 20-140 Annealed

H* 860 1640 55 170 70 Annealed

850 1640 52 200 70 Annealed

*longitudinal

;

+ transverse

22
1200

03 Q_

X'

I

o

z

LU

cc

\—

CO

Q

_l

m

>-

LU

CO

z

LU

1100

1000

900

800

700

600

500

0

Annealed 316LN

4K

0.05 0.10 0.15 0.20 0.25 0.30

GRAIN SIZE, d

Figure 1. Effect of grain size on the tensile yield strength at 4 K of two

316LN alloys. Using Eq. (2), for alloy B, = 839 MPa, m = 855

MPa; for alloy D, = 762 MPa, m = 570 MPa.

23
»

Figure 2. Effect of nitrogen content on the tensile yield strength (normalized to d = 0.50 ym) at ^ K of 31 6-type alloys. Data best fit the

relation Oy = 556 + 2450 [N] where SD for Oy = 64 MPa. Scatter

bands represent 2 SD.

24
TENSILE YIELD STRENGTH, MPa

Figure 3. The tensile yield strength versus the fracture toughness at ^ K for

316-type alloys. Scatter bands represent 2 SD.

25
X 10^

INCLUSION SPACING,

Figure The Quality Index Kj^(J)* 0 y

at 4 K versus the average inclusion

spacing (n~^^^) for 3'16-type alloys. Here n = the number of inclusions 0.5 ym) in 1 mm^. Data are corrected to Ni content

of 12 wt.%. Scatter bands represent 2 SD.

26
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STRENGTH AND TOUGHNESS OF AISI 30^ AND 316 AT ^ K

N. J. Simon and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

Structural design of superconducting magnets in fusion energy devices

requires reliable property data at 4 K. Nitrogen-strengthened AlSl 304 and

316 stainless steels are considered to be the best currently available low

temperature structural alloys on the basis of their f abricability and their

potential to meet the U.S. fusion research goals of combined 1000-MPa yield

strength and 200-MPa*m^ fracture toughness at 4 K. This paper presents

provisional equations for the yield strength of 304- and 316-type alloys at

4 K as a function of N content and grain size. Provisional equations for the

4-K fracture toughness, Kj^(J), of these alloys are presented, in which the

effects of yield strength, Ni and Mn content, and inclusion spacing are expressed quantitatively. The equations result from regression analyses of a matrix of NBS measurements that includes more extensive numerical data on alloying, refining, and processing parameters than have previously been

available .

1 . Introduction

Superconducting magnets for future fusion energy reactors will require

cryogenic structural alloys with improved combinations of yield strength

and fracture toughness at 4 K. Recently we have made considerable progress

in understanding how quantitative parameters of alloying, refining, and

processing can be specified for AISI 300 series stainless steels to achieve

the strength/toughness goals of the U.S. fusion program [1].

The structural alloy is required to withstand large magnetic forces in a superconducting magnet system of restricted dimensions, and thus, high

strength is needed. In addition, because most structures are complex and have

regions of high stress concentration, good fracture toughness is required.

Our research has focused on various AISI types of austenitic steels, especially 304 and 316. Metallurgical parameters affecting low temperature

mechanical properties have been studied using a series of laboratory heats.

Collaboration with recent Japanese austenitic steel development programs has

permitted mechanical characterization at 4 K of some of their new steels.

Compilation and critical review of physical and mechanical property data at

low temperatures have also been carried out.

*Work supported by the U.S. Department of Energy; not subject to copyright.

27
We have found that the strength at 4 K of austenitic steels depends

quite heavily on N content: each 0.1 wt.% N raises the tensile yield

strength (Oy, 0.2% offset) by approximately 300 MPa [2]. Additions of other

elements, such as C, Ni, Mn, and Cr , contribute much less significantly to the

low temperature strength. The yield strength is also dependent on grain

size: an increase of grain size from 20 to lOO ym results in a decrease of Oy

of about 90 MPa [3]. Thus, one can affect the strength through alloying and"^

thermomechanical processing.

Toughness is adversely affected by increased strength because less en- ergy is dissipated by plastic deformation during the void growth typical of

ductile fracture. We have shown that the elastic stress intensity factor

Kjq(J) [obtained from measurement of the elastic/plastic critical J-integral

(Jjq) and calculated from Kjq(J) = where E is Young’s modulus] is

linearly inversely proportional to o at 4 K [^1]. Toughness also is adversely affected by the presence of inclusions [3»5], which act as initiation

sites for microvoid formation. The addition of some alloying elements, such

as Ni, has previously been shown to increase the low temperature toughness of

austenitic steels when strength and inclusion content are held constant [6].

In comparing heat-to-heat strength and toughness at 4 K, we found a large amount of scatter in the data, about ±1005^. Such large scatter in

mechanical property data is typical at 4 K: a much smaller variation is

usual at room temperature [7]. To analyze this variability, we assembled a more extensive set of 4-K strength and fracture toughness data than has previously been available, combining our recent test data for the two highstrength cryogenic alloys, 304LN and 316LN, with earlier NBS measurements. We

determined quantitatively the effects Of grain size and N content on strength

and the effects of strength, alloying, and inclusions on toughness. Our

analyses resulted in predictive relations that considerably reduced the scatter of the data. Alloy design (in terms of strength and toughness) is now

feasible, within statistical limits of uncertainty. Of course, at this time

only research data are generally available. Before large-scale use of one of

these alloys in a low temperature structural application, additional heatby-fheat qualification testing is imperative.

2. Procedure

2.1. Data matrices

A series of specimens from both laboratory and commercial heats of AISI

304- and AISI 31 6-type alloys were tested. The alloys were prepared in the

United States, Japan, Europe, and the U.S.S.R. [3,4]. Measurements were made

of fracture toughness [Kj^(J)] (TL orientation), tensile yield strength, and

other tensile properties at 4 K. Grain size, hardness, and chemical composition were also determined and inclusion counts were obtained, as described by

R. P. Reed et al. [3]. Only measurements made at NBS were used to make up the

data matrices, because previous experience has indicated that lab-to-lab

variations in test results are large enough to obscure dependence of cryogenic properties upon parameters.

28
There are 25 complete rows in the Oy x Kjq(J) matrix for AISI 304-type

alloys and 19 complete rows in the corresponding matrix for AISI 31 6-type

alloys. The additional tensile data without corresponding Kj^(j) measurements were used in the regression analyses of the tensile yield strength.

Tables I and II show the data used in the regression analyses and the reference symbols. Tables III and IV present the complete chemical analyses and

hardness of the alloys. Nitrogen content ranges from 0.03 to 0.29 vit.% for

AISI 304-type alloys and from 0.03 to 0.17 wt./S for AISI 3l6-type alloys.

Nickel content varies from 7.2 to 14.9 wt.^K for AISI 304-type alloys, and

from 10.5 to 13.9 ^t.% for AISI 316-type alloys. Manganese content ranges

from 1.0 to 10.22 vit.% for the AISI 304-type alloys. Variation of Mn and Ni

content outside the AISI 304 specifications permitted an evaluation of the

effects of these elements; also, AISI 304-type alloys have recently been modified for cryogenic service by addition of Mn and Ni. Grain size varies from

19 to 385 urn in the AISI 304 matrix and from 11 to I 38 ym in the AISI 316

matrix. Inclusion spacing (defined as n~^^^ where n = the number of inclusions per mm^) ranges from 0.041 to 0.12 mm for the AISI 304 matrix and from

0.044 to 0.11 mm for the AISI 316 matrix. Hardness (Rg) was not used in the

analysis, but was found to vary from 72 to 90 for the AISI 304-type specimens

and from 79 to 100 for the AISI 31 6-type specimens. The K^^(J) values obtained from the relationship Kt^(J) = [Jt^*E/( 1-v^)]^ where v is Poisson's

ratio, were adjusted in accord with the formulation Kj^(J) = . All

data points of the matrices were weighted equally in the regression analyses

described below.

2.2. Analysis of tensile yield strength

Enough data were available to assess quantitatively the effects of grain

size and nitrogen content upon 4-K yield strength of both AISI 304- and AISI

316-type alloys. Earlier studies cited above have shown that these are the

most significant variables. In accord with the well-known Hall-Petch grain

size effect on yield strength and the higher temperature results of

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

^ °o

^ ®N

^ ®d

where [N] is wt.^ N, d is the grain size (ym), and is the yield strength

in the absence of N with infinite grain size. Quantification of the

strengthening effects of solid-solution alloying elements such as Ni would

require a more extensive data set, but are expected to be much less. The

contribution from C is expected to be about one-half that of N [2], but it is

best to hold C content low (-0.03 wt.^) to minimize sensitization during

welding

.

2 . 3 . Analysis of fracture toughness

The dependence of Kj^(J) upon yield strength at 4 K has been observed

previously in our laboratory [8,16], and a recent study of the effect of Ni on toughness of AISI 304-type alloys showed a strong dependence [6]. Manganese is

expected to have a similar, but weaker, effect. To assess the relative

importance of inclusion spacing with respect to the other two variables, the

following equation was fitted to the AISI 304 data:

29
Kic(J) = kg + kp Oy + k^ ([Ni] + 0.5 [Mn]) + k^^ (2)

where [Ni] represents wt.^ Ni, [Mn] ^t.% Mn, and kQ is a constant. Another

way of handling the effect of inclusion spacing is to to normalize Kj^(j) as follows:

Kj^(j)/n"‘''^2 = kQ* + kp’ Oy + kj^ ([Ni] + 0.5 [Mn]). (3)

This equation also was fitted to the 304 data. Since Mn content of the AISI

316-type alloys varied only slightly, the [Mn] term was dropped from

Equations (2) and (3) for the analyses of these alloys.

3. Results and discussion

3.1. Tensile yield strength

The equation for 4-K tensile yield strength of AISI 304-type alloys

resulting from a regression on a set of 30 measurements is:

Oy (MPa) = 215 + 3190 [N] + 1093 (4)

The standard deviation (SD) of the fit to the data is 84 MPa; the SD*s of

the three coefficients are 62, 230, and 383, respectively. For AISI 316-type

alloys, the comparable equation resulting from regression on a set of 26

measurements is:

Oy (MPa) = 381 + 2776 [N] + 811 d"’'^^. (5)

The SD of the fit to the data is 55 MPa; the SD*s of the three coefficients

are 34, 193i and 199. Figures 1 and 2 show the measured Oy versus the Oy

predicted by Equations (4) and (5) for both alloys.

The dependence of Oy on N in Equations (4) and (5) is slightly less than

that reported by Reed and Simon [2] from a different set of AISI 304-type

alloy data (not restricted to NBS measurements). The higher value of for

AISI 316-type alloys reflects the strengthening contribution of Mo. It is

not clear why we obtain a better data fit (lower SD) for the AISI 316-type

alloys. Many of the AISI 316 alloys are from commercial heats (17 out of

26); more of the AISI 304-type alloys are from laboratory heats (27 of 30).

Perhaps these data reflect better quality control of the producers in the

larger commercial heats.

3.2. Fracture toughness

The fit of Equation (2) to the 304 data gives:

Kj^(J) (MPa-m^^2) = 209 - 0.226 Oy + 14.4 ([Ni] + 0.5 [Mn])

+ 294 n"‘'/2 (6)

30
where the SD of the fit to the data is 31 MPa*m'''^^ and the SD’s of the

coefficients are 40, 0.026, 2.9 and 316. For the 316 data, the analogous

result is:

Kj^(J) (MPa-m''/2) = 130 - q. 338 Oy + 20.2 [Ni] + 2252 n‘

• 1/2 (7)

The SD of the fit to the data is 31 MPa*m^'^^, and the SD's of the coefficients are 89, 0.059, 5.2, and 310. Figures 3 and 4 show the measured K-j-^CJ) versus Kj^(J) predicted from Equations (6) and (7). Equations (6) and (7)

imply that the fracture toughness of the AISl 316 alloys is more sensitive to

all major variables than that of the AISl 304 alloys. The explanation of

this awaits more detailed studies of the effects of metallurgical variables

on the J-integral fracture toughness at low temperatures. In the set of

measurements analyzed here, the AISl 304-type laboratory heats generally have

a higher level of inclusions than the commercial heats. However, for the

AISl 316 -type alloys the situation is reversed: the laboratory heats are cleaner .

The SD's of the fits of both alloy types to Equation (2) were considered

to be small enough (31 MPa) that the equations could give useful predictions

for alloy design. However, there are significant differences in the coefficients for AISl 304- and 31 6-type alloys. In contrast, the coefficients

resulting from fits of Equation (3) to both alloy series were remarkably

similar. The SD's of these fits to Equation (3) were higher than for the

fits to Equation (2), so these equations are not recommended for predictive

purposes at present, but are set forth here for their heuristic value.

304: Kj^(J)/n ^ 5555 _ ( 8 )

The SD of the fit [to Kj^(J)] is 76 MPa-m'*'^^; the SD's of the coefficients

are 1324, 0.92, and 101 .

316 : Kj^(J)/n '’^2 ^ 5024 - 4.80 0 218 [Ni] (9)

The SD of the fit [to Kj^(J)] is 38 MPa*m^^^; the SD's of the coefficients

are 1577, 1 .0, and 92.

_ Compared with a fit of Kj^(J) on 0 alone, the additions of both the

n term and the [Ni] or ([Ni] + 0.5(lMn]) term were effective in reducing

the SD of the fits, except for Equation (6), in which the addition of the

n term does not affect the SD. This term is presented for comparison

with Equation (7). It is difficult to measure inclusion spacing

unambiguously [ 3 ], but inclusions do have a pronounced effect upon the 4-K

fracture toughness of the AISl 316-type alloys. Further inclusion studies

should be carried out on commercial heats of AISl 304-type alloys.

3.3 Alloy design

Equations (4) through (7) can be used to set provisional design values

for yield strength and fracture toughness. For example, to achieve the U.S.

research goal of a = 1000 MPa [1] for a 304LHN(31 6LHN) with [N] = 0.20 ut.%,

a grain size of 551162) ym is required. [The uncertainty in grain size is

31
± 25(4^) ym, assuming that the the fractional error in grain size is given by

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

other terms of the equation. The fractional error of each term is obtained

by dividing the SD of the coefficient by the value of the coefficient.]

Since the control of grain size to 55 ym is ^difficult for most mills when

rolling thick-section (50 - 120 mm) plate, this implies that AISI 31 6-type

alloys would be a sounder choice to achieve such high strength.

Assuming further that [Ni] = 12 wt.^ and [Mn] = 2 wt.?, achievement of

1 /P the U.S. research goal of = 200 MPa*m would require a minimum inclusion spacing of 0.10(0.073) mm for AISI 304LHN(31 6LHN) . However, the uncertainty is too large: ± 0.11(0.056) mm. This indicates that additional

studies are required to characterize the effect of inclusions on the fracture

toughness of austenitic steels at 4 K. Further studies to relate composition

and other variables to the weldability of these alloys are also necessary [17].

4 . Summary

Matrices of NBS measurements of Oy and Kj^(J) at K that include nu- merical data on alloying, refining, and processing variables were analyzed

using linear regression methods. The AISI 304- and 31 6-type alloys tested

were obtained from a variety of sources and included both commercial and

laboratory heats.

Tensile yield strength: Provisional equations were developed for

estimating Oy at 4 K from [N] and

Fracture toughness: Provisional equations for estimating Kj^(J) at 4 K

from Oy, [Ni], [Mn], and n~^^^ were determined. In addition, the predictions

of the quantity Kj^(J)/n~^^^ from Oy and [Ni] were found to be similar,

within the statistical uncertainty ,'^for both AISI 304- and 316-type alloys.

The equations presented are expected be useful for alloy design and for

tentative estimates of attainable fracture toughness and tensile yield

strength at 4 K for austenitic stainless steels.

5. Acknowledgments

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

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

R. P. Walsh for conducting many of the tensile and fracture toughness tests

and to P. T. Purtscher for the inclusion analysis. L. L. Scull also greatly

assisted us by determining interstitial contents in most of the alloys.

AISI 316-type alloys were supplied by AMAX Materials Research Center

(P. J. Grobner), and Nippon Kokan Steel, Technical Research Center (C. Ouchi)

kindly supplied a series of laboratory heats. Other alloys were provided by

Lawrence Livermore National Laboratory (E. N. C. Dalder), Oak Ridge National

Laboratory, and Kernforschungszentrum Karlsruhe, Institut fiir Technische

Physik (A. Nyilas). The laboratory heats of AISI 304-type alloys were suppled by U.S. Steel, ARMCO , Carpenter Technology, and the E. 0. Paton

Institute of Electrowelding, Kiev, U.S.S.R. (K. Yushchenko).

32
6. References

[ 1 ]

[ 2 ]

[3]

[iJ]

[5]

[ 6 ]

[7]

[ 8 ]

[9]

[ 10 ]

[ 11 ]

[ 12 ]

H. I. Morlenry, in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures - VIII, Internal Report 85-3025, Ed.

R. P. Reed (U.S. National Bureau of Standards, Boulder, Colorado, 1985)

p. 21.

R. P. Reed and N. J. Simon, in: Advances in Cryogenic Engineering - Materials, vol . 30, Eds. A. F. Clark and R. P. Reed (Plenum Press, New

York, 198^) p. 127.

R. P. Reed, N. J. Simon, P. T. Purtscher, and R. L. Tobler, Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures - IX, Internal Report, Ed. R. P. Reed (U.S. National Bureau of Standards,

Boulder, Colorado, 1986) p. 15.

R. L. Tobler, D. T. Read, and R. P. Reed, in: ASTM STP 7^13 (American

Society Testing Materials, Philadelphia, Pennsylvania, 1982), p. 250.

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

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

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

R. P. Reed and P. T. Purtscher, Advances in Cryogenic Engineering - Materials, vol. 32, Eds. R. P. Reed and A. F. Clark (Plenum Press, New York,

1986) p. 43; in press.

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

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

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

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

Energy Applications at Low Temperatures - III, Internal Report 80-1627,

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

1980) p. 15.

R. P. Reed, Test results of an AISI 304-type alloy. National Bureau of

Standards; unpublished data.

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

Energy Applications at Low Temperatures - II, Internal Report 79-1609,

Eds. F. R. Fickett and R. P. Reed (U.S. National Bureau of Standards,

Boulder, Colorado, 1979) p. 79.

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

Energy Applications at Low Temperatures - I, Internal Report 78-884,

Eds. F. R. Fickett and R. P. Reed (U.S. National Bureau of Standards,

Boulder, Colorado, 1978) p. 91.

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

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

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

33
[13] R. L. Tobler, R. P. Mikesell, R. L. Durcholz, C. W. Fowlkes, and

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

Superconducting Machinery, Internal Report 7^-359, (U.S. National

Bureau of Standards, Boulder, Colorado, 197^) p. 182.

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

Energy Applications at Low Temperatures - IV, Internal Report 81-1645,

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

Boulder, Colorado, 1981) p. 147.

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

[16] R. L. Tobler, T. A. Siewert, and H. I. McHenry, in: Material Studies

for Magnetic Fusion Energy - IX, Internal Report, Ed. R. P. Reed (U.S.

National Bureau of Standards, Boulder, Colorado, 1986) p. 239 .

[17] E. N. C. Dalder and M. C. Juhas, Nucl. Eng. Des. (1986) to be published.

34
Table II. ^ Matrix, AISI 31 6-type Alloys

Alloy M?;

N.

wt . %

d,

pm

^Ic ’ 1 / p MPa*m '

Ni,

wt . %

Mn,

wt . % mm

A [9] 935 0.16 63 — 13.7 1 .26 —

B [9] 1076 0. 1 7 13 219 11.4 1 .40 0. 1

1

942 0.17 70 — 11.4 1 .40 —

924 0.17 91 324 11.4 1 .40 0. 1

1

917 0.17 1 29 313 11.4 1 .40 0.1 1

C C9] 834 0.10 31 350 13.9 1 .84 0.095

834 0.10 31 390 13.9 1 .84 0.095

834 0.10 31 350 13.9 1 .84 0.095

D [9] 1056 0.12 — 150 10.9 1 .01 0.082

935 0.12 1

1

180 10.9 1 .01 0.082

885 0.12 20 190 10.9 1 .01 0.082

869 0. 12 30 — 10.9 1 .01 —

842 0.12 50 -- 10.9 1 .01 —

E [10] 1010 0.17 105 200 13.9 1.58 0.057

1010 0.17 105 214 13.9 1.58 0.057

F [10,12] 71

1

0.03 65 266 13.5 1 .86 0.054

711 0.03 65 254 13.5 1 .86 0.054

71

1

0.03 65 243 13.5 1 .86 0.054

663 0.03 67 — 13.5 1 .86 —

656 0.03 67 — 13.5 1 .86 —

G [13,14] 550 0.04 64 400 12.2 1 .55 0.092

690 0.04 22 — 12.2 1 .55 —

661 0.04 28 — 12.2 1.55 —

640 0.04 32 — 12.2 1 .55 —

553 0.04 49 — 12.2 1.55 —

540 0.04 55 — 12.2 1.55 —

562 0.04 58 — 12.2 1.55

535 0.04 61 — 12.2 1.55 —

535 0.04 81 — 12.2 1.55 —

530 0.04 118 — 12.2 1 .55 —

505 0.04 138 -

12.2 1.55 --

H [9] 859 0.15 70 159 10.5 1.10 0.044

859 0.15 70 181 10.5 1 .10 0.044

854 0.15 70 197 10.5 1 .10 0.052

854 0.15 70 203 10.5 1.10 0.052

*n = the number of inclusions per mm^.

35
Table I. °y x Matrix, AISI 304-type Alloys

Alloy

N,

Wt . 1o

d,

pm

^Ic ’

1 /p

Ni,

wt . % Mn,

wt. ?

n"*! /2

,

* mm

A [8] 880 0.16 85 218 9.91 1 .50 0.063

1190 0.29 98 118 9.97 1 .42 0.054

530 0.04 90 320 10.1 1 .57 0.041

740 0.1 4 105 250 9.9 1 .50 0,057

1290 0.26 90 118 10.0 1 .42 0.055

B [9] 620 0.11 57 150 8.68 1 .02 0.044

760 0.11 57 230 8.62 5.81 0.041

880 0.19 55 110 7.81 1.19 0.047

1050 0.21 53 150 7.83 5.79 0.057

C [9] 970 0,20 150 110 7.36 4.17 0.047

1090 0.21 130 1 40 7.15 10.22 0.050

1040 0.21 125 160 9.09 4.33 0.054

980 0.19 100 140 7.02 7.5 0.054

D [9] 350 0.03 210 330 9.72 1 .28 0.044

E [10] 770 0.13 57 210 8.60 1 .74 0.123

F [10] 1340 0.28 55 91 8.44 1 .83 0.107

G [6] 950 0.28 200 143 8.9 4.2 0.063

460 0.09 325 273 8.7 3.7 0.057

1000 0.26 385 204 12,8 5.5 0.066

720 0.14 340 244 11.5 3.9 0.057

1160 0.28 225 239 14.9 5.2 0.070

870 0.20 320 254 14.7 4.4 0 . 070

H [9] 720 0.12 75 202 10.8 1 .02 0.074

680 0.12 80 212 10.8 1 .02 0.074

817 0.12 19 — 10.8 1 .02 —

784 0.12 23 — 10.8 1 .02 —

719 0.12 1 61 — 10.8 1 .02 —

720 0.12 161 — 10.8 1 -02 —

697 0.12 50 21 4 10.8 1 .02 —

I [11] 794 0.12 120 21 4 9.49 1 .88 0.101

p

* n = the number of inclusions per mm

36
Table III. Chemical Composition and Hardness of AISI 304 -type Alloys

Chemical Composition, wt.^ Hardness,

Alloy Cr Ni Mn C N P S Si Cu Mo

A 18.0 9.91 1 .50 0.067 0.16 0.022 0.017 0.57 0.20 83

20.3 9.97 1 .42 0.030 0.29 0.024 0.016 0.55 0.20 — 89

18.3 10.1 1 .57 0.089 0.04 0.021 0.019 0.50 0.21 — 79

18.0 9.9 1 .50 0.037 0.14 0.022 0.017 0.57 0.20 — 79

20.3 10.0 1 .42 0.085 0.26 0.024 0.016 0.55 0.20 — 90

B 18.2 8.68 1.02 0.025 0.11 0.022 0.013 0.53 0.30 — 82

18.1 8.62 5.81 0.021 0.11 0.023 0.014 0.64 0.30 — 84

19.4 7.81 1.19 0.025 0.19 0.023 0.015 0.60 0.28 — 88

19.5 7.83 5.79 0.033 0.21 0.024 0.014 0.61 0.31 — 90

C 19.4 7.36 4.17 0.025 0.20 0.020 0.01

1

0.50 0.25 0.31 82

19.4 7.15 10.22 0.023 0.21 0.019 0.01

1

0.44 0.25 0.31 76

21 .3 9.09 4.33 0.025 0.21 0.020 0.01

1

0.48 0.25 0.31 82

19.6 7.02 7.50 0.023 0.19 0.019 0.01

1

0.47 0.24 0.30 84

D 18.4 9.72 1.28 0.051 0.03 0.030 0.016 0.48 — 0.32 75

E 18.8 8.60 1 .74 0.024 0.13 0.038 0.003 0.51 — 0.38 79

F 20.2 8.44 1 .83 0.059 0.28 0.025 0.021 0.50 0.60 0.59 —

G 18.4 8.9 4.2 0.018 0.28 0.014 0.005 0.28 — 82

18.7 8.7 3.7 0.014 0.09 0.013 0.005 0.23 — — 72

20.5 12.8 5.5 0.024 0.265 0.015 0.004 0.34 — — 81

19.7 11 .5 3.9 0.012 0.1 41 0.013 0.005 0.27 — — 72

20.8 1 4.9 5.2 0.028 0.277 0.015 0.004 0.41 — — 82

20.8 14.7 4.4 0.016 0.197 0.004 0.014 0.33 — — 76

H 18.0 10.8 1 .02 0.018 0.12 0.002 0.003 0.28 — — 84-77

I 18.6 9.49 1 .88 0.048 0.12 0.019 0.024 0.38 0.07 0.52 79

37
Table IV. Chemical Composition and Hardness of AISI 316-type Alloys

Alloy

Cr Ni Mn

Chemical composition, wt.

Mo C N P

%

S Si

Hardness

A 16.7 13.7 1 .26 2.7 0.032 0.16 0.01 0.01 0.41 89

B 18.5 11.4 1 .40 2.5 0.008 0.17 0.002 0.001 0.40 89

C 16.8 13.9 1 .84 2.2 0.028 0.10 0.022 0.009 0.33 88

D 18.6 10.9 1 .01 2.2 0.020 0.12 0.002 0.003 0.29 96-102

E 17.4 13.9 1 .58 2.5 0.016 0.16 0.021 0.021 0.48 83

F 17.2 13.5 1 .86 2.3 0.057 0.03 0.024 0.019 0.58 79

G 17.3 12.2 1 .55 2.2 0.051 0.03 0.026 0.014 0.68 86

H 18.0 10.5 1.10 2.1 0.01 6 0.1 4 0.020 0.015 0.54 86

38
MEASURED

TENSILE

YIELD

STRENGTH,

Oy,

MPa

Figure 1 . Measured tensile yield strength at 4 K of AISI 30^-type alloys

versus tensile yield strength calculated from Equation (A). The

scatterband indicates ± 2 SD's. Keys to symbols representing

test specimens are given in Tables I - IV.

39
MEASURED

TENSILE

YIELD

STRENGTH,

Oy,

MPa

500 700 900 1100

PREDICTED TENSILE YIELD STRENGTH, Oy, MPa

Figure 2. Measured tensile yield strength at 4 K of AISI 3l6“type alloys

versus tensile yield strength calculated from Equation (5). The

scatterband indicates ± 2 SD’s. Keys to symbols representing

test specimens are given in Tables I - IV.

40
MEASURED

FRACTURE

TOUGHNESS,

K|c(J),

MPaFigure 3. Measured fracture toughness at M K of AISI 304“type alloys versus fracture toughness calculated from Equation (6). The scatterband

indicates ± 2 SD’s. Keys to symbols representing test specimens are given in Tables I - IV.

41
0 100 200 300 400 500

PREDICTED FRACTURE TOUGHNESS, K|c (J), MPa-m’'^"

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

versus fracture toughness calculated from Equation (7). The

scatterband indicates ± 2 SD’s. Keys to symbols representing

test specimens are given in Tables I - IV.

42
"4

^ •* •

'

• til

^/^''* fraot^d AiJd

‘ *

' Centj4^^o^;»atel3|4i!^'^oi%e^

-i

:r

Vi

^\4|

*

.*

•

NEAR-TMRESHOLD FATIGUE CRACK GROVTH BEHAVIOR

OF AISI 316 STAINLESS STEEL*

R. L. Tobler

Fracture and Deformation Division

Center for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

The near-threshold fatigue behavior of an AISI 316 alloy was characterized using a newly developed, fully automatic fatigue test apparatus.

Significant differences in the near-threshold behavior at temperatures of 295

and M K are observed. At 295 K, where the operationally defined threshold at

10“^° m/cycle is insensitive d stress ratio and strongly affected by crack

closure, the effective thresh.old stress intensity factor (AK'pj^)Qff is about

4.65 MPa*m'/^ at R = 0.1 and n = 0.3. At 4 K, the threshold is higher, crack

closure is less pronounced, and there is a stress ratio dependency; (AK'pj^)gff

is 5.1 MPa‘m^^^ at R = 0.3 and 6.1 MPa*m^^^ at R = 0.1. There is also a significant difference in the form of the da/dN-versus-AK curves on log-log

coordinates: at 4 K the curve has the expected sigmoidal shape, but at 295 K

the trend is linear over the region of da/dN from 10~® to 10“^° m/cycle.

Other results suggest that the near-threshold measuremehts of a 6.4-mm-thick

specimen of this alloy are insensitive to cyclic test frequency below 40 Hz.

INTRODUCTION

Near-threshold fatigue crack growth rates (da/dN-versus-AK curves) and

threshold stress intensity factors (AK-pj^) are relevant mechanical parameters

for superconducting magnet design since they lead to the prediction of safe

operating conditions below which failures from fatigue crack growth can not occur. Owing to measurement difficulties and costs, however, the existing

data base at 4 K is still quite limited. Liaw and coworkers ^ designed a fully automatic fatigue test apparatus and reported threshold data for

precipitation hardenable Fe-29Ni-1 4Cr-1 . 2Mo austenitic stainless steel and its

welds at 297, 77, and 4 K. Tobler and Cheng designed a similar apparatus and

presented data for AISI 304L, 304LN, and 316 stainless steels at 4 K. The

present paper considers the AISI 316 alloy in greater detail, contrasting the

behaviors at 295 and 4 K and exploring the effects of some mechanical test

variables.

*To be published in A^’vrnces in Cryogenic Engineering—Materials, volume 32.

43
(J,

^1/

I

{J.

MATERIAL

The tested naterial was a 50 . S-mm- thick plate of AISI 316 austenitic

stainless steel in the a:mealed condition. T.he chemical composition in wt.$

is ?e-*~.T50n-T5.-?Ni-0.05^T- I . S6Mn-2 . S^lMo-O ..0214P-0 . 01 9S-0.58Si-0.030N.

Tensile, fracture toughness, and midrange fatigue crack growth data were previously.® Selected tensile properties are listed in Table 1,

th Young's moduli for a sim.ilar grade of steel.®

NTAL rROOETYRI

iver

fatig-ue test apparatus used in this study includes a 100-kN

machine and cryostat with minicomputer, programmable

generator, programmable digital oscilloscope, and liquid

m. The procedure is fully automated, allowing continuous

load adjustments to be performed in the absence of a Tests were conducted at -0 Hz and lasted about days with

ration. T>-pioally, *3 million load cycles are needed to

old measurement and to aocumulate rate measurements ranging

m ovole.

.ne t . --zn-tnio; ' >

•? »- this study is sho«T: in Fig.

and the load cell and gage

1 .

'Vice.

ection-versus-load curves they were continuously

'llosoo”'e t’^aces "^or tests

Ct Si.. ^ n

^ -H . ^S3KS

I
Here P^ax ^min maximum and minimum applied fatigue loads, and Pq]_

is the closure load as identified in Fig. 2. Additional details of the test

apparatus and procedures are published elsewhere.**

RESULTS

The da/dN-versus-AK curves for AISI 316 are presented in Fig. 3* The

da/dN values are plotted versus applied AK values without correction for

crack closure. The rates of 10~® to 10“^° m/cycle are two orders of

magnitude lower than rates previously published for this alloy, and at the

threshold level m/cycle), they correspond to one Angstrom per load

cycle or less than one lattice spacing per load cycle.

At 295 K (Fig. 3A), the rates at R = 0.1 and R = 0.3 are nearly

indistinguishable, showing no effect of stress ratio. The data trend is

linear on log-log coordinates, and the rates do not exhibit a knee or approach an asymptote at the near-threshold level. Read and Reed^ tested the same heat of AISI 316 at AK from 25 to 50 MPa*m^/^. Their data, representing

25.4-mm-thick specimens at 20 Hz, are in excellent agreement with and

complement the present results.

The results at 4 K (Fig. 3B) differ from those at 295 K in that there is

a significant effect of R ratio. The rates at R = 0.3 are higher than at

R = 0.1, and the threshold stress intensity factor is larger at R = 0.1. At

this temperature, the da/dN curves do exhibit the expected sigmoidal shape,

and there is approximate agreement with the results for 25 . 4-mm-thick

spec imens .

®

The effect of test temperature at R = 0.1 is also shown in Fig. 3C. The

linear and sigmoidal trends at 295 and 4 K are distinct and cross twice. The

AK-p^ value at 4 K is slightly higher than that at 295 K, and when crack

closure corrections are applied, the temperature effect is more pronounced.

This can be seen by comparison of the ( AK'p}^)gfp values in Table 2. It is

evident from these data that reducing the temperature from 295 to 4 K

increases the threshold fatigue resistance.

Figure 4 presents the effect of frequency variations at 4 K on the crack

growth rate at AK of 20 and 42 MPa*m^/^. For these tests the fatigue

apparatus was programmed for constant AK, and replicate crack growth rate

measurements were performed. Frequency was varied by two orders of magnitude

from 0.4 to 40 Hz. When the range of experimental error and some effects of

crack closure and variability are taken into account, these data show that

there is no measurable effect of test frequency on the crack growth rates.

DISCUSSION

At 295 K, the fatigue crack growth rates for AISI 316 appear to follow a linear trend for at least four orders of magnitude in da/dN, and no asymptote

is approached at the threshold level of 10"^° m/cycle. This has several

implications. First, it implies that no true threshold is exhibited in this

45
range. Hence, the operational definition of AK-pj^ at 10“^° m/cycle is

inadequate for this alloy/ temperature combination, and a new definition must

be sought. Second, it points to the necessity of performing additional

measurements at still lower growth rates. Note that an average growth rate

of 10“‘° m/cycle corresponds to 1 A/cycle, while the lattice spacing for

austenite is about 3.6 A. Rates of crack growth less than a lattice

parameter per cycle are possible if the crack g,rowth occurs nonuniformly

,

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

In these tests, crack closure became more significant as the cracks

propagated and the loads were reduced. As shown in Fig. 2, the crack closure

was much more significant at 295 K than at 4 K. The plasticity-induced

closure mechanism can be used to explain this observation. According to this

mechanism, closure is attributed to residual stresses left in the wake of the

advancing crack tip plastic zone. ® This closure mechanism would contribute

greater closure effects at 295 K where the yield strength is lower. The fact

that the threshold AK values at R = 0.1 and R = 0.3 are larger at 4 K than at

295 K can be explained as an intrinsic effect of the increased lattice

resistance to dislocation movement at cryogenic temperatures.

Frequency effects on da/dN have been observed in AISI 316 at elevated

test temperatures because of creep interactions with fatigue®. At low

temperatures creep is less consequential and cyclic frequency is expected to

have little influence on fatigue crack growth rates. At T = 4 K, however, an exception is conceivable, since local heating can occur in the crack-tip

plastic zone due to plastic strain energy dissipation. The amount of heat

generated must decrease at lower cycle frequencies and at lower AK. Since

Fig. 4 shows no measurable variation in rates at AK - 20 or 42 MPa*m^/^

between 0.4 and 40 Hz, and since the plastic zone size is significantly

smaller at lower AK, we tentatively conclude that the AK-pj^ and (AK'ppj)gff

parameters for AISI 316 are not frequency dependent below 40 Hz.

SUMMARY AND CONCLUSIONS

This study of the near-threshold fatigue crack growth rate behavior

of annealed AISI 316 stainless steel at 295 and 4 K gives the following

results

:

1 . At 295 K, the da/dN-versus-AK curves on log-log coordinates follow a linear trend for four orders of magnitude. The fact that the data do not

approach an asymptote at 10“^° m/cycle leads to the conclusion that still

lower rate measurements are necessary to reach a true fatigue threshold, and

that the conventional operational definition of AK-pp^ is not satisfactory for

this material and temperature combination.

2. At 4 K, the da/dN-versus-AK curve is sigmoidal, and the operational definition of AK-pj^ is satisfactory at this temperature.

3. The near-threshold behavior of AISI 316 is stress ratio dependent

at 4 K, but not at 295 K, for R ratios between 0.1 and 0.3.

46
. The threshold stress intensity factor for this material is temperature dependent, increasing as temperature is ref iced from 295 to 4 K.

ACKNOWLEDGMENTS

This project was supported by the Office of Fusion Energy (DOE),

Dr. V. Der , project monitor.

REFERENCES

1. P. K. Liaw, W. A. Logsdon, and M. H. Attaar, in: "Austenitic Steels at

Low Temperatures, " R. P. Reed and T. Horiuchi, Eds., Plenum Press, New

York (1 983), pp. 171-185.

2. R. S. Williams, P. K. Liaw, M. G. Peck, and T. R. Leax, Eng. Fract.

Mech. 18:953-964 (1983).

3. P. K. Liaw, W. A. Logsdon, and M. H. Attaar, in: "Fatigue at Low

Temperatures, " ASTM STP 859, R. I. Stephens, Ed., American Society for

Testing and Materials, Philadelphia, Pennsylvania (1985),

pp. 173-190.

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

growth rate measurements at liquid helium temperature, Int . J . Fat . 7:191-197 (1985).

5. D. T. Read and R. P. Reed, in: "Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures—II," NBSIR 79-1609, National

Bureau of Standards, Boulder, Colorado (1979), pp. 81-122.

6. H. M. Ledbetter, J. Appl. Phys . 52:1587-1589 (1981).

7. R. J. Bucci, in: "Fatigue Crack Growth Measurement and Data Analysis,"

ASTM STP 738, S. J. Hudak, Jr. and R. J. Bucci, Eds., American Society

for Testing and Materials, Philadelphia', Pennsylvania (1981), pp. 5-28.

8. S. Suresh and R. 0. Ritchie, in: "Fatigue Crack Growth Thresholds

Concepts," D. Davidson and S. Suresh, Eds., Met. Soc. AIME (1984), pp.

227 - 262 .

9. L. A. James, At. Energy Rev. 14:37-85 (1976).

47
Table 1

.

Mechanical properties and grain size of the test material

.

Test Yield Ultimate Red. of Young's ASTM

Temperature Strength, Oy Strength, Ou Elongation Area Modulus, E Grain Size

(K) (MPa) (MPa) (%). (%) (GPa) No.

295 228 576 56 73 195 5.0

71 1 1301 57 208

Table 2. Nominal and effective threshold stress intensity

factors for AISI 316 stainless steel.

Test Temperature

(K)

Stress Ratio, R (MPa'^m‘i'/2)

(AKTh)gff

(MPa*m' ^2)

295 0.1 7.6 4.6

0.3 7.6 4.7

J

4 0.1 8.5 6.1 0.3 6.4 5.1

Fig. 1. Compact test specimen used in this study.

48
c

o

E

E

00

1—

H

o

o

c

o

X>

_0)

M—

0> Q

Load, 0.2 kN/division

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

<V Load, 0.4 kN/division

(B) T = 4K,AK = 8.5MPa.m‘/2

Fig. 2. Deflection-versus-load curves for AISI 316 test specimens at the

termination of threshold fatigue tests: (A) T = 295 K; (B) T = K.

49
Fatigue

Crack

Growth

Rate,

da/dN,

m/cycle ABC

Fig. 3. Fatigue crack growth rate data for AISI 316 stainless steel:

(A) at 295 K, (B) at 4 K, and (C) comparison at R = 0.1

50
Fatigue

Crack

Growth

Rate,

m/cycle

Fig. Effect of cyclic test frequency on 4-K fatigue crack growth rates at

two intermediate AK values.

51

I y Utd;^ strength

^lijr.y ' at"".*! 'X. '-auljstit^i

*S> Infiuervoing trto ^»'4«s

'/

f^^ture

* ^

'.D*Pnf> V awT. ;' t/* .'Vfu?a^ sj^

f.‘ '.. • Andepcnd^j^t, of to© yie|^ ©fe

"W^e;

yoy WOPO©?. Iffi ©uiKUlWItSA'^Xe#,^^

^6^ur*o tough04?a©:'

l*)j

THE TOUGHNESS OF AUSTENITIC STAINLESS STEELS AT 4 K*

P. T. Purtscher and "R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

The fracture toughness values of twenty-four different compositions of

austenitic stainless steels were measured in liquid helium in accordance with

the ASTM E 8l3“8l method. Three independent factors were found to contribute

to the energy required for the initiation of ductile fracture: yield strength

at 4 K, substitutional alloying (Ni, Mn, and Mo content), and nonmetallic

inclusions. Yield strength and inclusions have long been recognized as influencing the toughness of structural alloys and are considered in existing

models for fracture toughness. The effects of alloying on the toughness are

not directly considered in these models. Experimental results show that

substitutional alloying elements can influence the deformation and fracture

behavior of these steels in the toughness test in a manner that is not

independent of the yield strength or inclusion content.

Key words: austenitic stainless steels; cryogenic properties of materials;

ductile fracture; fracture mechanisms; fracture toughness;

mechanical properties

*Work supported by the U.S. Department of Energy.

53
INTRODUCTION

The construction of high-field superconducting magnets requires

structural alloys of high strength and good toughness at 4 K. Austenitic

stainless steels 304LN and 3 I 6 LN, strengthened by N additions (an interstitial alloying element), have been used in this application. In the past

ten years considerable data for the mechanical properties of austenitic

stainless steels at 4 K have been published. An empirical relationship

between yield strength and fracture toughness at 4 K was first reported by

Tobler et al. :

Fracture toughness = 500 - 0.3 x yield strength (1)

where yield strength is expressed in MPa and fracture toughness in MPa/m.

This equation represents the combination of strength and toughness found in

interstitially strengthened 304-type steels.

Superconducting magnet systems now under construction have been designed

with existing 304LN and 31

6

LN alloys. New concepts for magnet designs

require alloys with better properties. The Japanese Atomic Energy Research

Institute (JAERI) has proposed a specification for the strength and toughness

of structural alloys that cannot be obtained with 304LN or 3 I 6 LN. In the

United States, design goals^ for structural alloys Include a slightly lower

yield strength level than the Japanese counterpart, 1000 MPa rather than

1200 MPa. The fracture toughness requirement is the same. The U.S. design

goals also include strength and toughness requirements for welds. For

critical areas in the structure where the stresses are the highest, a welding

process is needed that will produce welds whose mechanical properties match

those of the base metal in all welding positions. For less critical locations in the structures, the toughness requirement is 255^ lower than that for

the base metal, and a minimum of two other welding processes are targeted.

The different design goals are shown in Figure 1 relative to the properties

of 304 -type stainless steels.^

The relationship between strength and toughness in materials has been a subject of research for years and is well-known in steels,^ AS, alloys,^ and

Ti alloys^ that fail in a ductile manner. The ductile fracture process

involves the initiation and growth of voids in a plastically deformed region

of the material. Typically, the voids initiate at nonmetallic inclusions

leaving small ligaments between initiation sites. The voids grow until they

coalesce with other voids or a free surface. The extent of void growth is

related to the strength and <iuctility of the matrix.

Current knowledge of the strength-toughness relationship in austenitic

stainless steels at 4 K is summarized in a recent review article by Morris

and Dalder.° They concluded that there are only two ways to improve the

strength-toughness relationship, that is, to increase toughness without

lowering the strength. The first way is to reduce the inclusion content.

With fewer void initiation sites, there will be fewer voids and the growth

Stage should be longer before the voids coalesce. This method of improving

54
the strength-toughness relationship is related to processing methods (such as deoxidation practice, ingot size used, and hot rolling) during steel production. The second way is to distribute the deformation over a larger area.

If the deformation is concentrated in a small region ahead of the crack tip,

the measured fracture toughness should be relatively low. If the size of the

plastic zone at the crack tip increases, "the toughness should Increase. This

method is related to the slip characteristics of the material, namely the

interaction of moving dislocations with obstacles to their motion. Cross

slip is one type of interaction that is expected to play a major role.

Dislocation interactions are complex and difficult to study. Fundamental theories have employed dislocation interactions successfully to

explain the effect of grain size,"^’^ second-phase particles, ^ and alloying

additions^ on strength. No similar fundamental theories exist for ductile

fracture.

Understanding ductile fracture is hindered by its varied definitions:

energy absorbed during an impact or static test at a given temperature,

transition temperature for ductile to brittle behavior, and blunt machined

vs. sharp fatigue precracked notches. All these factors influence ductile

fracture and can produce contradictory results. The effect of grain size on

Kic for -AISI 43^0 is opposite to the effect of grain size on the Charpy V-notch

energy for the same steel. Work by Lee et al.^^ explains clearly why blunt

and sharp notches have different effects.

Continuum mechanics models for ductile fracture attempt to relate

fracture toughness to tensile properties. Hirth and Froes^ summarized the

many different theoretical relationships found in the literature and

identified four factors that should be considered: ( 1 ) the relationship of

toughness to the plastic zone size and to the crack opening displacement;

( 2 ) the relationship of crack opening displacement to the spacing of

void-initiating inclusions; ( 3 ) the inverse relationship of fracture

toughness to strength, owing to the presence of plastic instabilities that

produce failure along characteristic slip paths; and ( 4 ) significant variations in the surface area created during fracture due to crack tortuosity and

crack branching. A more recent model for ductile fracture has been proposed

by Ritchie and Thompson”' ^ that incorporates a different term, the microroughness, M. The M term reflects the local strain to failure from measurements of the shape of the voids produced during fracture. In our discussion

we shall use this model to analyze our data.

The purpose of this study is to assess the role of alloying and inclusions in changing the toughness of austenitic stainless steels. The alloys

were chosen carefully so that the effect of substitutional alloying elements

(Ni, Mn, and Mo) could be studied; other . variables that influence the

strength and toughness were kept to a minimum. This information is particularly valuable in designing alloys that will meet specifications for service

at 4 K with cost effective use of alloying additions.

55
MATERIALS

Alloys with Nickel Additions

Eight completely austenitic alloys, in the form of 25~mm-thick plates, were used to evaluate the effect of Ni content; their chemical compositions

are listed in Table I. The as-received plates were annealed at 1060'^C for

1 h and then water quenched. The inclusion content of each of the eight

alloys was determined to be sure that there were no significant differences

among the eight alloys. A polished cross section of a tensile specimen

from each alloy was examined in the scanning electron microscope (SEM) at a magnification of 2000. The number of inclusions greater than 0.5 pm in

diameter in a given area (either 0.3^ or 0.17 mm^) was used as a means of

characterization. Most inclusions were the MnSi type, but there were also

CrMn-rich inclusions present. This technique of character izing inclusions is

similar to that used in another study its results were given in terms of

where n was the number of inclusions with a diameter larger than 0.5 ym

in a 0.1 mm2 area. For the eight alloys of this study, all counts were normalized to the same 0.1 mm^ area, and the parameter was calculated

to be 0.15 to 0.24.

Alloys with Manganese Additions

Two groups of alloys were prepared in which Mn content was the major

test variable. The first group (69-78) was tested in 1982; results are reported elsewhere. The results are analyzed differently here and compared

with those of the second group. Two types of inclusions were found in this

first Mn group with the SEM: MnS and AlSiCa-rich inclusions. Figure 2 shows

a typical distribution of these inclusions. There were between 30 and 60

inclusions per 0.1 mm^^ which translate to values of 0.13 to 0.1

8

for n

^

In addition to these inclusions, analysis by scanning transmission

electron microscopy (STEM) revealed extremely fine inclusions (0.1 pm in

diameter), rich in Ti and Ail (see Figure 3). These smaller inclusions were

found most often in the 1? Mn alloys and, to a lesser degree, in the 6 % Mn

alloys.

The four alloys from the second group (297, 304, 300, and 301), supplied

as 25-mm-thicl< plate, were austenitic with 1 to 4? second phase, either delta

ferrite or sigma. Their compositions are listed in Table I. The alloys were annealed at 1160°C for 6 h and then water quenched. Figure 4 shows the

microstructure of one of the alloys with second phase. The inclusions were examined in the SEM as described above. The n parameter values were

0.15 to 1.17. Many types of inclusions were present: MnS, MnSi, AilMn-rich,

and AilSi-rich inclusions.

Alloys with Molybdenum Additions

Molybdenum effects were studied on specimens from two different alloys

where the composition differed only in Mo (see Table I). Plates, 25 mm thick, were supplied in the hot -rolled condition and were annealed at temperatures

between 1000 and 1150°C for varying lengths of time in order to recrystallize

56
the alloys and grow the grains to different siz

measured: varied from 0.22 to 0.30. The

containing Ail, Mg, Si, S, Ca, and Ti.

es. The inclusion density was inclusions were complex.

EXPERIMENTAL PROCEDURES

Tensile Testing

Tensile tests were conducted in liquid helium with a cryostat described

elsewhere. Round specimens, 6-mm gauge diameter and 25-mm gauge length and

machined transverse to the rolling direction, were used for most of the

tests. Flat, dog-bone-type specimens with a gauge length of 25 mm were used

in one case (noted in Table II). Load vs. elongation curves were obtained by means of a commercial load cell and strain gauge extensometers . Yield

strength at 0.255 strain, ultimate strength, percent elongation, and percent

reduction in area were measured. Estimated strain sensitivity is

10 ME at 4 K.

Fracture Toughness Testing

Fracture toughness specimens, about 25 mm thick, were machined in a T-L

orientation with respect to the rolled plate. The specimens were tested at

4 K according to the ASTM E 8l3“8l test method, the single specimen

technique.^ The elastic unloading compliance method and a clip gauge^^

developed especially for cryogenic service were used to monitor crack growth

during the test. Scatter in the value measured at 4 K resulted in

uncertainties of ±155^.

RESULTS

The tensile properties and the fracture toughness values for all the

steels are shown in Table II. Yield strength (Oy) vs. Kjq(J) fracture

toughness is plotted in Figure 5 along with the scatter band of the previous

study of 304-type steels with different interstitial contents. The expected

general trend of decreasing toughness with increasing strength is seen, but

there is considerable scatter from the previous data on 304-type steels.

Another way to analyze the data is to consider the influence of the

substitutional alloying elements within each group of alloys, and thus, to

isolate this influence.

Nickel

The effect of Ni was presented in another paper^^ along with the effect

of N, but the main results are reiterated here and expanded. The effect of

interstitial elements on the yield strength of these alloys is shown in

Figure 6. Complete stress-strain curves for the four alloys with the same N content but different Ni contents (5.6 to 14.955) are shown in Figure 7. T'r.«> initial strain hardening characteristics are similar. The strain to faiiurdoes vary, but it is not simply related to the Ni content.

57
The load drops that occur with strain hardening are related to adiabatic

heating of the specimen,

a

conversion of the mechanical work done on the

specimen by plastic deformation to heat. Owing to the low specific heat and

thermal conductivity of austenitic stainless steels, the specimen temperature

can rise, lowering the yield strength and promoting plastic deformation.

Serrated yielding is common in austenitic stainless steels tested at 4 K. The size of the load drops in alloys 413“^16 appears to increase with Ni

content, from 180 MPa for alloy to 460 MPa for alloy 4l6. This

difference could be related to the temperature dependence of the yield

strength. For alloy 413, almost all the plastic deformation was associated

with the load drops, that is, temperature rises. The other three alloys

showed appreciable plastic strain at 4 K.

Fracture toughness is presented as a function of Ni content in Figure 8.

Five of the alloys had nearly the same interstitial content and yield

strength. The measured toughness values of the other three alloys were normalized to a yield strength of 1030 MPa, the average of the first five

alloys, to minimize the effect of the different N contents. After accounting

for the different yield strengths, Kj^(j) increased by about 18 MPa/m for

each percent increase in Ni content.

The load vs. load-line displacement curves for alloys 413 through 41 6,

from which the fracture toughness values were determined, are shown in

Figure 9. The point on each curve at which crack growth began is noted with

a circle. These four alloys vary only in Ni content. Nickel affects the

shape of the curve prior to crack initiation, promoting deformation prior to

crack initiation under conditions of severe constraint. Note the sharp load

drops that occur after is passed. These load drops are always associated

with an audible "pop" that indicates that the crack is growing discontinuously

.

Audible sounds did not accompany the load drops during the tensile tests shown

in Figure 7.

The four alloys with high nitrogen but varying Ni contents, 413, 41 4,

415, and 4l6, show a distinct transition in the fracture morphology. A low

magnification view of the fracture surface in the region of the specimens where

the fracture toughness is measured (Figure 10) shows this distinct transition

from faceted fracture to dimpled fracture between 9 and ^ 2 % Ni.

In some areas of the fracture surface of alloy 41 3, like that shown in

Figure 11, the size of a facet is equal to the grain size, and two or three

sets of parallel lines are faintly visible on the facet. The angles between

the different sets of lines are approximately equal to the angles between

different slip directions in the fee crystal. The plane of the facet has

been identified by back-reflection Laue patterns to be a {111} type plane.

Another region where the facets are extremely small and stepped in appearance

is shown in Figure 12. Here, the {111} planes are not favorably oriented with

respect to the tensile load, and the crack follows small steps between {111}

planes. The nonmetallic inclusion does not appear to have affected the

fracture process.

58
A higher magnification view of the fracture surface that is characteristic of alloys ^415 and ^16 with the dimpled appearance is seen In Figure 13.

For these alloys, the inclusions played a significant role in the fracture

process. The voids initiated at the inclusions and grew until they coalesced

with surrounding voids.

Strain-induced martensite did form in some of these alloys during

mechanical testing. Measurements of martensite content were made and are reported in the earlier work. From those results, it was determined that

martensite did not have a direct effect on the toughness values.

To summarize the results, a linear change in toughness was found as a function of Ni content after the effect of yield strength variation was eliminated. This change in Kj^^(j) was associated with alloys exhibiting

brittle, faceted fracture as well as ductile, dimpled fracture behavior.

The inclusion content of all the alloys is similar; their influence on the

variation in fracture toughness is minimial.

Manganese

The results of testing the first group of alloys with Mn as a major

alloying element (69 through 78) were reported previously . ^ ^ The results

(see Figure 14) showed a systematic increase in toughness with Mn content at

three different N levels. However, the toughness values of all the Mn alloys

were lower than expected from the results of the Fe-l8Cr-10Ni study; the

difference in toughness was noted in the original paper but not examined

thoroughly.

In this study, the fracture surfaces of the toughness specimens were examined in detail. The fracture appearance of the alloys offers an explanation for the unexpectedly low toughness of the Mn alloys. Figure 15

shows the fracture surface of alloy 69, which contained '\% Mn. The fracture

is relatively smooth at low magnification and appears to have very fine

dimples, 1 to 2 pm in diameter, over most of the surface. The fine dimples

initiated at the Ti- and AS,-rich inclusions; the density of these inclusions

decreased with increasing Mn content (1 to in the alloys.

_

The fracture surface of alloy 72 containing Mn (Figure 16) appears

rougher than that of alloy 69 (Figure 15), and its dimples are considerably

larger, about 20 pm in diameter. The dimples have formed around larger MnS

inclusions rather than fine Ti- and A2,-rich inclusions. The difference in

toughness between the \% and 655 Mn alloys (Figure 14) appears to be a function of both dimple size and Mn content.

The fracture toughness values for the second group of Mn-alloyed steels

are shown in Figure 17 in two different 'ways; normalized to a yield strength

of 1030 MPa and normalized for yield strength plus a 7/5 Ni content, according

to the trend noted in Figure 8 for Ni alloying. The toughness averages for

three of the alloys show a trend of increasing Kj^(J) with increasing Mn

content that is approximately one-third that found for Ni. The toughness

of the fourth alloy did not follow this trend, but normalizing the toughness

59
to account for the difference in Ni content brought it closer to the trend

line. These results show considerable scatter, possibly due to the 1 to

of second phase present.

In the second group of alloys, a transition in fracture morphology,

similar to that seen for Ni alloying in Figure 10, was observed and is shown

in Figure 18. The transition in fracture appearance is associated with the

Mn content: the alloy with Mn, 297, has a faceted fracture surface, and

the alloys with higher Mn content, 30^ and 300, have dimpled fracture

surfaces. Alloy 301, with Mn but a slightly higher Ni content, is not

shown, but it had a dimpled appearance similar to that of alloys 304 and 300.

To summarize, Mn appears to increase the 4-K toughness, but the magnitude of the effect is difficult to judge. The results of the two groups are shown together in Figure 19. The two different trends are significant

because the steeper trend is thought to reflect an effect of dimple size and

Mn content, while the lower trend is a function of Mn content alone.

Increasing the Mn or Ni content of alloy 297 does change the fracture

appearance significantly, indicating an important role for substitutional

alloying elements in the deformation and fracture processes.

Molybdenum

The measured yield strengths of the Mo-alloyed specimens are shown

in Figure 20. The measurements from specimens with similar grain sizes are

noted. Solid-solution strengthening from 2.2^ Mo increased the 4-K yield

strength of those specimens with a similar grain size by about 80 MPa.

The toughness values of the Mo-alloyed group are plotted in Figure 21

,

again with the measured values normalized to a strength of 1030 MPa. The

2.256 Mo addiction increased _the adjusted fracture toughness from 105 MPa/m up

to 165 MPa/m (about 25 MPa/m per wt. % of Mo added).

The fracture appearance of the Mo-alloyed group was completely dimpled

and did not change with grain size. However, the dimple sizes were different

for the alloys with and without Mo. Figure 22 shows alloy 80, without Mo.

The dimples are extremely fine, approximately 2 ym in diameter. No inclusions were associated with the fine dimples. Figure 23 shows that the

dimples in alloy 79 (with Mo) are generally much larger and contain inclusions.

General Observations

Certain pertinent observations were common to all specimens tested:

1 . Fractography of the broken specimens . No stretch zones were observed in the region of the fracture surface between the fatigue

precrack and the J-test-related features. This is characteristic

of austenitic stainless steels that are broken at cryogenic

temperatures .21

60
2. Nature of the ductile crack growth . Load drops were observed in

the load vs. load-line displacement curves from which the toughness

is calculated. Load drops were associated with an audible sound,

indicating discontinuous crack growth.

DISCUSSION

The results show a significant effect of substitutional alloying

elements on fracture appearance and mechanical properties, particularly

fracture toughness. The role of substitutional alloying elements in the

fracture process has not been directly incorporated in any existing theory.

To explain how Ni, Mn, and Mo can affect the fracture toughness, a simple

mechanism is proposed for fracture of annealed materials:

There are three basic stages in the development of the fracture

toughness value that we measure in the ASTM E 8l 3 procedure: Stage 1 - The

growth of a plastic zone ahead of the sharp fatigue precrack. This stage occurs in all cases. Stage 2 - The creation of a free surface on the

microscopic scale within the plastic zone. For ductile fracture, this

typically involves the decohesion of the matrix-inclusion interface.

Stage 3

- The growth of that microscopic free surface into the macroscopic

crack, which is defined as crack initiation. For ductile fracture, this

stage is associated with the gradual linking of voids.

The size of crack that is defined as the point of initiation is not

clearly specified by the ASTM E 8l 3 standard. The initial standard issued in

1981 requires a crack growth vs. resistance curve to be generated, and is

defined as the intersection of the resistance curve, extrapolated to zero

growth, and the blunting line. This technique provides a reliable, consistent method of evaluating the toughness, but it does not always provide a simple correlation to such metallurgical factors as inclusion size and

spacing, dimple size and spacing on the fracture surface, tensile properties,

interfacial energies for decohesion of inclusion-matrix interfaces, and

stacking-fault energies.

Influence of Alloying

The effect of substitutional alloying elements appears to be continuous

through the brittle-to-ductile transition (Figure 7 and 18), so it seems reasonable to consider stage 1 , the growth of the plastic zone, as the only step

in the fracture mechanism where Ni, Mn, and Mo can affect both ductile and

brittle fracture equally.

The growth of the plastic zone in the fee structure is related to

dislocation motion (glide) and the ability of dislocations to cross slip,

effectively reducing the stress concentration at the tip of dislocation

pileups. Ahlers^^ described cross slip in fee metals as a thermally

activated, strain-rate-sensitive process by which extended dislocations on

one glide plane must constrict and come together to form a perfect disloca61
tion before they can change to another glide plane. Extended dislocations

are composed of two partial dislocations with a spatial separation that is

related to the stacking fault energy. A concise description of the effect of

stacking fault energy on mechanical properties can be found in several

sources . 23,24 For our case, it is enough to say that cross slip is more difficult in an alloy with a low stacking fault energy, and that it can produce significant effects on the deformation behavior. In the toughness

test at 4 K, temperature and strain rate are not variable, and the effect of

composition on the stacking fault energy is the parameter most likely to

influence cross slip and, hence, the plastic zone formation (stage 1).

A factor that could affect cross slip but was not mentioned by Ahlers is

stress state. Figure 7 shows that increasing Ni content has no significant

effect on the engineering stress-strain curves at 4 K where a uniaxial state

of stress exists. However, Figure 9 shows a dramatic effect of Ni on the

deformation behavior of these same alloys where a triaxial state of stress is

present.

The transition in fracture morphology seen in Figures 10 and 19 can be

interpreted in terms of the fracture mechanisms and the effect of composition

on the cross slip characteristics. With the high-strength, but lower-alloyed

steels (413, 414, and 297), brittle facets formed because the slip is

concentrated on the {111}-type planes. The addition of alloying elements (Ni

in 415, Mn in 304, and Ni in 301) relieved the concentrated slip, promoting

dimple formation.

Qualitatively, the cross slip model can explain why steels with different

Ni, Mn, and Mo contents have different toughness values. Unfortunately, it

is difficult to measure stacking fault energy. Schramm and Reed^^ reported a relationship between stacking fault energy (SFE) and content of major

alloying elements for commercial-grade austenitic Fe-Cr-Ni, Fe-Mn-Ni, and

Fe-Cr-Ni-Mn steels:

SFE = -53 + 6.2(5S Ni) + Cr) + 3-2(5^ Mn) + 9.3(5^ Mo) (2)

where SFE is expressed in mJ/m *wt./5. Other researchers have reported significantly different equations relating chemical composition to stacking

fault energy.

Equation 2 predicts high values of SFE, but we feel that it gives a good prediction of the relative effect of the major alloying elements, that

is to say, Ni is 2 times as effective as Mn, Mo is 3 times as effective as

Mn, and although Cr has an effect, it is considerably less. These relative

effects correlate with the trends for the effect of alloying elements on toughness that are reported here.

Influence of Dimple Size

The test results for the first Mn group and the Mo group appear to be a function of both alloy content (stage 1) and dimple size (stage 2 and 3). At

present, there is no way to separate the effects of the two factors. One

possible solution would be to use a continuum mechanics approach to

62
understanding fracture. Ritchie and Thompson's model^ 3 fo^* ductile fracture

incorporates both i';icrostructural and fractographic observations into the

calculation:

Jlc = (of/3) Un C(h/W)2/3fp] il], (3)

where o

0

0

fyu

h = W =

flow strength i (Oy + a^)/2,

yield strength,

ultimate strength,

average height of dimple,

average width of dimple,

volume fraction of inclusions associated with dimples,

character istic distance, taken to be 5 times the average spacing

of void-initiating inclusions

The key element in their model is that the shape of the dimple (h/W) and the

inclusion content are simply related to the local strain to failure. They

assume a stress-modified critical strain criterion where J = when the

local strain to cause failure is exceeded over a characteristic distance that

is comparable to the inclusion spacing.

This model can be applied to fracture surfaces that are relatively

uniform in appearance and where there are distinct differences in dimple

morphology. In our study, the best example of contrasting dimple shape is

found in the first Mn group, with extremely fine dimples in the \% Mn alloys

(Figure 15) and large dimples in the 6? Mn alloys (Figure 16).

To apply the model, certain assumptions need to be made concerning the

shapes of the dimples and volume fraction of inclusions. The fractographs in

Figures 15 and 16 show that the large dimples, about 20 ym in diameter,

appear to be approximately half as deep as they are wide (h/W = 0.5), whereas

the fine dimples, about 2 ym in diameter, appear to be relatively shallow

compared to their width. The h/W parameter is assumed to be 0.25, and the

character istic distance is assumed to be 5 times the average dimple diameter.

The f^ of the 0.1 ym diameter inclusions is estimated from the STEM foils to

be 10 to 10 The f of the larger inclusions in the 6? Mn alloy was estimated to be 0.015.

With these measurements and assumptions and the measured tensile

properties, the calculated values of Jj^ for alloy 69 (1^ Mn) were 21 and

29 kJ/m^ for the two different values of fp; the measured value was 109 kJ/m^.

For alloy 72 (6? Mn) the calculated was 65 kJ/m^; the measured value was

260 kJ/m^. The model predicts a lower toughness than that measured, but it

does predict the correct trend of increasing toughness with increasing dimple

size—the toughness ratios for the two alloys, [(21 or 29)/65] calculated and

[ 109 / 260 ] measured, are approximately the same.

Another way to test the model is to compare two alloys with very similar

dimple shapes. Take, for example, alloy 69 from the first Mn group and alloy

80 from the Mo group. We can see that the shapes of the dimples are similar

(compare Figures 15 and 22), the volume fractions of void-initiating

inclusions should be reasonably close, and the character istic distances

should be the same in both cases. With these assumptions, the model predicts

63
that toughness is a function of the flow stress alone. For these two alloys,

the flow stress is approximately the same, 1170 and 1133 MPa. The calculated

toughness values for the two are similar, between 21 and 29 kJ/m^. However,

the measured values found in Table II were considerably different, 109 for

alloy 69 and 216 kJ/m^ for alloy 80. We conclude that this model does not

adequately describe these test results.

If the fracture toughness values of the two alloys are normalized to the same yield strength, 677 MPa, the toughness of alloy 80 is still 212 MPaVm,

and for alloy 69, the K-]-^(j) value is 132 MPaVm. Adjusting the toughness to

account for the difference in the Ni content 0 0.8 for 80 and 8.7 for 69)

increases the value for alloy 69 to 168 MPaVm. Now, the comparison between

the fracture toughness values of the two alloys is improved.

A possible explanation for the difference in toughness is that a different type of inclusion is responsible for void nucleation, stage 2 in our simple model. When the toughness values of different alloys within one group are compared, the type of inclusion does not vary. In alloy 69, the

fine inclusions are rich in Ti and Al. The fine inclusions assumed to exist

in alloy 80 have not been identified, but because the alloys were made by

different suppliers, it seems likely that the type of inclusion is different.

These two examples illustrate specific shortcomings of the model

presented by Ritchie and Thompson. The major problem is determining the

proper values of the variables: flow stress, dimple shape, and characteristic distance. In addition, there is no method of incorporating the effect of

alloying elements into the model if the elements do not affect the tensile

properties, for instance, the effect of Ni on the deformation behavior in the

toughness test as compared to the tensile test.

The logarithmic term for local strain to failure provides a good,

physical model of the process, but only drastic changes in the fracture

features (h/W), types of fracture (dimpled or quasi -cleavage) or shape

(such as shown here for alloys 69 and 72) are quantitatively important in

evaluating the toughness.

The characteristic distance is a variable that can be changed to make

the data fit the model. In Ritchie and Thompson's presentation of the model,

5 to 6 times the average major inclusion spacing was an appropriate value for

ductile fracture of ASTM A533B-1 steel. Lee et al. have shown that 1 to 2

times the inclusion spacing is appropriate to use for high-strength

AISI 43^0 steel. An analysis of data for high-strength ki alloys (see

Appendix A) indicates that S,* would also be 1 to 3 times the average spacing

of major void-initiating inclusions for these alloys.

Another problem with the existing model is that it does not take into

account the different types of inclusions that initiate voids. The tensile

ductility of ferritic steels has been shown to vary significantly as a function of inclusion type and shape.

A

similar effect of inclusion type

could explain why two alloys with similar dimple shape and flow strengths

have different fracture toughness values.

64
SUMMARY

The results reaffirm previous results that show toughness decreases with

increasing strength. However, at a constant strength level, Ni, and possibly

Mn and Mo, significantly increase toughness. . This increased toughness is

linked to a difference in dislocation interaations and is evidenced by

increased plasticity for a triaxial stress state and the transition in

fracture morphology. For brittle, faceted fracture, dislocation pileups

produce decohesion along slip planes. With higher Ni content, Mn content, or both, dimpled fracture occurs. For the austenitic stainless steels tested,

we suggest that the alloy additions increase the stacking fault energy and

make cross slip easier. More cross slip increases plasticity and reduces the

stress concentration at the head of dislocation pileups. Then, the preferred

site for decohesion would be the interface between inclusions and the matrix.

Another factor affecting toughness appears to be the size of dimples

produced in ductile fracture. The Ritchie-Thompson model has correctly

predicted the trend of decreasing toughness with decreasing dimple size, but

the prediction is not quantitative. Problems occur when attempting to

correlate the model quantitatively with the test results.

ACKNOWLEDGMENTS

The Office of Fusion Energy, Dr. Victor Der, Project Monitor, sponsored

this research. We owe special thanks to Robert Walsh and Ralph Tobler who

performed most of the mechanical testing. We would also like to thank

David Gorni and Chris McCowan for their help in reviewing the document and

suggesting ways to improve it.

REFERENCES

1. R. L. Tobler, D. T. Read, R. P. Reed: in Fracture Mechanics : Thirteenth

Conference , ASTM STP 7 ^ 3 , R. Roberts, eds., American Society for Testing

and Materials, "Philadelphia, 1981, p. 350.

2. H. I. McHenry: in Materials Studies for Magnetic Fusion Energy Applica - tions at Low Temperatures - VIII , NBSIR 85-3025, R. P. Reed, ed..

National Bureau of Standards, Boulder, Colorado, 1985, p. 121.

3. J. M. Krafft: Appl. Mat. Res., vol. 3, 1964, p. 88.

4. G. T. Hahn and A. R. Rosenfield: Metall. Trans. A, vol. 6A, 1975,

p. 653.

5. J. P. Hirth and F. H. Froes

:

Metall. Trans. A, vol. 8A, 1977 , p. 1165.

6

.

J. W. Morris, Jr. and E. N. C. Dalder: J. Met., vol • 37, no. 10, 1985,

P. 24 •

7. J. P. Hirth: Metall. Trans. , vol. 3, 1972, p. 3047.

8. J. C. M. Li and Y. T. Chou

:

Metall. Trans., vol. 1, 1970 , P. 1145.

65
9.

C. W. Corti, P. Cotterill, and G. Fitzpatrick: Int. Mater. Rev.,

vol. 19, June 197^, p. 77.

10. J. C. Fisher, Acta Metall., vol. 2, 195^, p. 9.

11. R. 0. Ritchie and R. M. Horn: Metall. Trans. A, vol. 9A, 1978, p. 331.

12. S. Lee, L. Majno, and R. J. Asaro: Metall. Trans. A, vol. 16A, 1985,

p. 1633.

13 . R. 0. Ritchie and A. W. Thompson: Metall. Trans. A, vol. I 6 A, 1985,

p. 233 .

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

S. Shimamoto: in Advances in Cryogenic Engineering - Materials , vol. 30 , Plenum Press, New York, 1984, p. 137.

15 . R. L. Tobler and R. P. Reed: in Advances in Cryogenic Engineering - Materials , vol. 28, Plenum Press, New York, 1982, p. 83 .

16 . R. P. Reed: in Advances in Cryogenic Engineering - Materials , vol. 7,

Plenum Press, New York, 1961, p. 448.

17 . ASTM E 813 - 8 I, in "Standard Test Method for Jj^, A Measure of Fracture

Toughness, 1981 Annual Book of ASTM Standards , Part 10, American Society

for Testing and Materials Philadelphia, 1981, p. 763 .

18 . R, L. Tobler and J. A. Shepic: J. Test. Eval. July 1985, p. 299.

19 . R. P. Reed and P. T. Purtscher: in Advances in Cryogenic Engineering - Materials , vol. 32, Plenum Press, New York, 1986, p. 43.

20. D. T. Read and R. P. Reed: in Advances in Cryogenic Engineering - Materials , vol. 26, Plenum Press, New York, 1980, p. 91.

21. P. T. Purtscher: submitted for publication to J. Test. Eval.

22. M. Ahlers: Metall. Trans., vol. 1, 1970, p. 2415.

23 . H. Suzuki: Sci. Rep. Res. Inst. Tohoku Univ., vol. 4A, 1952, p. 455.

24. B. Burton: Acta Metall., vol. 30, 1982, p. 905.

25 . R. E. Schramm and R. P. Reed: Metall. Trans., vol. 6A, 1975, p. 1345.

26 . R. L. Widner, W. K. Dominik, and A. J. Jenkins: SAE paper 780784,

Automot. Eng., vol. 86, no. 9, 1978, p. 74.

66
Table I. Chemical Composition of Steels (wt.^)

Steel Cr Ni Mn C N Mo n-l/2*

for Ni Alloying

^413 18.8 5.6 4.1 0.020 0.256 — 0.15

415 18.4 8.9 4.2 0.018 0.218 — 0.20

41 4 20.5 12.8 5.5 0.024 0.265 — 0.21

416 20.8 1 4.9 5.2 0.028 0.277 — 0 . 18 - 0.22

520 19.8 5.8 3.2 0.01 4 0.260 — 0.24

739 18.7 8.7 3.7 0.014 0.093 -- 0.18

740 19.7 11.5 3.9 0.012 0.1 41 — 0.18

714 20.8 14.7 4.4 0.016 0.197 — 0.22

1 St group for

Mn Alloying

69 18.2 8.7 1.0 0.025 0.1

1

0.32 0.1 4 70 18.1 8.7 2.0 0.025 0.11 0.31 —

71 18.2 8.7 4.0 0.025 0.1

1

0.31 —

72 18.1 8.6 5.8 0.021 0.1

1

0.33 0.13

73 19.4 7.8 1.2 0.025 0.19 0.32 0.15

74 19.4 7.9 2.0 0.022 0.19 0.32 —

75 19.3 7.8 3.9 0.024 0.19 0.31 —

76 19.5 7.8 5.8 0.033 0.21 0.31 0.18

77 18.7 8.2 4.0 0.026 0. 15 0.31 —

78 18.4 8.3 5.8 0.031 0.15 0.31 —

2nd group

for Mn Alloying

297 19.4 7.4 4.2 0.025 0.20 0.31 0.15

304 19.6 7.0 7.5 0.023 0.19 0.31 —

300 19.4 7.2 10.2 0.023 0.21 0.31 0.1 7 301 21 .3 9.1 4.3 0.025 0.21 0.30 0.1 6

Mo Alloying

79 18.6 18.9 1 .0 0.020 0.12 2.2 0.24-0.30

80 18.0 10.8 1 .0 0.018 0.17 0.0 0.23

*n = number of inclusions greater than 0.5 ym in diameter' in a 0 .

1

2 mm"^ area

67
Table II. Mechanical Properties of Steels

MPa

0u»

MPa

elong.

,

%

R.A.

,

%

^Ic»

kJ/m^

Kic(J),

MPa/m

Grain Size,

ym

for Ni Alloying

41 3 1075 1418 1

1

17 27 75 122

414 950 1650 42 54 99 143 165

415 1000 1580 32 — 201 204 200

416 1 1 60 1630 22 48 276 239 160

520 960 121 6 7 — 37 87 180

739 460 1 417 40 43 360 273 220

740 720 1370 42 50 288 244 250

741 870 1 425 29 48 312 254 207

1 St group for

Mn Alloying*

69 616 1 724 3^ 23 109 150 57

70 661 1616 29 22 163 184 55

71 722 1581 38 31 203 205 49

72 755 1506 44 34 260 232 57

73 883 1808 34 30 60 1 1

1

55

74 956 1524 22 20 64 115 55

75 986 1668 38 31 98 142 53

76 1053 1584 33 25 116 155 53

77 857 1654 38 31 188 197 53

78 901 1574 41 37 185 196 53

2nd group for

Mn Alloying

297-1 /2 880 1 700 47 45 71 /66 128/117 150

304-1/2 950 1530 41 45 65/112 1 1 6/152 100

300-1 /2 1020 1590 37 43 93/112 139/152 130

301 -1/2 1025 1580 45 46 99/1 32 1 43/165 125

Mo Alloying

79-1 1056 1733 36 49 103 146 7 79-2 935 1686 41 54 151 177 1

1

79-3 885 1618 41 58 169 187 20

80-1 677 1588 41 50 216 212 80

80-2 715 1620 42 51. 199 202 75

80-3 784 1654 42 55 23

80-4 817 1669 41 57

*flat, dog-bone-type specimens were used for this group of alloys

68
YIELD STRENGTH, Oy (MPa)

Figure 1 . The Japanese and United States design goals for strength and

toughness of austentic stainless steels at 4 K relative to the

properties of 304-type stainless steels.^

69
Figure 2. Appearance of nonmetallic inclusions in scanning electron

microscope. Size and distribution is representative of all

the steels included in this report

70
0.1 ym

Figure 3* TEM bright field and dark field micrographs showing the 0.1 ym

diameter inclusions found in alloys 69~78. Region A shows an inclusion, and region B is an area of the foil where an inclusion was before it fell out during electropolishing.

71
Figure

Figure 5.

0.2 mm

Microstructure of alloy 30^ with second phase in the

austenitic matrix.

ch

if)

LLI z

I

o

Z)

R

LU

a;O

<

cc

350

300

250

200

150

100

50

0

400 600 800 1000 1200 1400

YIELD STRENGTH, o, (MPa)

1 ^ I I I r V ;\

y—

N

\—

X

\—\/

Previous Results from

Fe-18Cr-10Ni‘

\ ^

A ^

\

X

VV

X

X

X

X Ni Alloying 413-416, 520, 739-741 Mn Alloying 69-78 O Mn Alloying 297-304 • Mo Alloying 79 A Mo Alloying 80

J ^ ^ I L

\

\ \

\ V

X

Fracture toughness vs. yield strength for all of the alloys

presented in this report.

72
Figure 6. Yield strength at M K as a function of interstitial alloying

elements.

73
Stress,

MPa

Stress,

MPa

Strain

Fig. 7a and b. Engineering stress vs. strain diagrams

of alloys ill 3 and 414 at 4 K.

74
Stress,

MPa

Stress,

MPa

Strain

Strain

(c)

(d)

Fig. 7c and d. Engineering stress vs. strain diagrams

of alloys 415 and 4l6 at 4 K.

75
350

03 CL

300

250

200

(/) to

o>

c: sz

OJO Z3o

150

a>

-4—1 o

03 100

50 -

0>-

0

• High-Nitrogen alloys

o Normalized for yield

strength of

1030 MPa

5 10 15 20

Nickel Content (wt. %)

Figure 8. Fracture toughness vs. Ni content 1 for alloys

413-416, 520, and 739-741 at 4 K.

76
413 414

2000 kN

O03O

4000 kN

J|q = 27 kJ/m2

5.6% Ni

Load-Line Displacement, u

Q_

TD

CDO J|c = 99 kJ/m2

8.9% Ni

Load-Line Displacement, u

Figure 9. Load vs. load-line displacement curves for alloys ^13, 414, 415,

and 416 at 4 K.

77
Figure 10. SEM fractographs of the fracture surface of alloys 413, 414, 415

and 41 6 that were created during the toughness test.

78
Figure 11. SEM fractographs from the fracture surface of alloy 413 showing

large facets that are nearly equal to the grain size.

79
5ym

I I

Figure 12. SEM fractograph from the fracture surface of alloy 413 showing

smaller facets in the vicinity of a nonmetallic inclusion.

20 ym

Figure 13. SEM fractograph from the fracture surface of alloy 41 6.

80
YIELD STRENGTH, Oy (MPa)

Figure 14. Fracture toughness vs. yield strength for alloys 69~78

(different Mn contents) at 4 K.

81
100 ym

20 ym

Figure 15. SEM fractograph from the fracture surface of alloy 69 with Mn.

82
100 ym

Figure 16. SEM fractograph from the fracture surface of alloy 72 with 655 Mn.

ra Qcn

C/5 LUI

O

Z)R

UJ

QC

ID

I

o

<

QC

LL

200 -

190 -

180

170

160

150 A

140

130 /

A ^ /

120 /

110

100 •

^—Transition from

/ brittle-to-ductile

90 — o \ A = compact tension

80 - O = bend bar

open symbols are corrected

70 A to a yield strength of 1030 MPa

60 solid symbols are corrected

^

^ L_J ^ 1 ^ L_J for strength and 7% Ni 01 234 56789 10 Manganese Content (wt.%)

Figure 17. Adjusted fracture toughness vs. Mn content for alloys

297, 300, 301 , and 304, at 4 K.

83
Figure 18. SEM fractographs of the fracture surface of alloys 297, 304, and

300 that were created during the toughness test.

84
Figure 19

03 Q_

C/D CO

LUZ

I

o

IDP

LU

OC

Z)

I

o

<

QC

Li200

1

-

190

180

1701-

160

150

140

130

120

110

100

90

80

70

60

Dimple

size changes—^ /

o r

/

7

/

/ /

V /

/

/

/

/

\

/ ^Transition from

/ brittle-to-ductile

./

/

/

41 0123456789 10

Manganese Content (wt.%)

. Adjusted fracture toughness vs. Mn content for both Mn-alloyed groups.

85
1100

1000

'c5' I

900

>% o

B 800

O)c0w_ 05^ 700 0

i>-

600

500

Molybdenum Content, (wt.%)

—^

X

X

0

0

0

X

X

— X =as measured —

0 =as measured with

^ ~ 20/jm grain size

0 12 3 4

Figure 20. Tensile yield strength vs. Mo

content for alloys 79 and 80

at 4 K.

1^

CO

CO

LUX

o

3

LU

CLD

O

<

CL

200

190 -

180 -

170 -

160

150

140

130

120

110

100

90

80

70

60

o /

/

y

^Dimple

/ size changes

O = corrected to a

strength of 1030 MPa

Figure 21 . Fracture toughness vs. Mo 1 1 1 ' content for alloys 79 and 80 0123

at 4 K. Molybdenum Content (wt.%)

86
Figure 22. SEM fractograph of the fracture surface of alloy 80, 0% Mo.

20 ym

Figure 23. SEM fractograph of the fracture surface of alloy 79, 2.2% Mo.

87
APPENDIX A. INFORMATION ON ALUMINUM ALLOYS

This information is from a paper by Hahn and Rosenfield. If the shape of

the dimples is assumed to be hemispherical, h/W is equal to 0.5. Combine

this assumption with the data in Table A-1 and one can see that the 2,^ term

is 1 to 3 times the average spacing of major inclusions, not 5 or 6 times, as described in the text.

Table A-1 . Data for High-Strength Aluminum Alloys

201 i1-T6 2024-T851 7079-T651

Oy (MPa) 436 443 502

Kjq (MPa/m) 19.5 20.5 27.4

Ufc (kJ/m^) 1.8 2.1 3.6

dp (ym) 8.4 7.4 9.4

Vf {%) 2.9 4.4 0.9

then 2,q (ym) is 12 22 10

dp = average spacing of major inclusions

v^ = volume fraction of major inclusions

2-^ = characteristic distance for ductile fracture

REFERENCE

A-1. G. T. Hahn and A. R. Rosenfield; Metall. Trans. A, vol. 6A, 1975,

p. 653.

88
^^ >? iritETCif :jms

. "f •'m\r l*JJ^ t • -!S '

-

H. ""‘€

i| 'j*^, -?r' ''_r'-'’* ,* 'X“ ^ ._ ip|)flfpri>:v -op-^ pli

t ' tty xitrani - . #tr«t<!f*<wf

.

tr-f" ^:«' » kBi^'

.M-. i V

le .purpetBt? pf M

t4ffr ^nn»tljQfi 0 Btr

KLiXa;

jpp •’ \ <^h.9 ,. C9>

:

Jffs.

' '«nl. r«ri' v.iv A •Moi.

'.tteohnoiWjki; Cnglneera^,; .^f'SB^'iiM!’* -to

ABSENCE OF STRETCH ZONES IN AUSTENITIC STAINLESS STEELS

FRACTURED AT CRYOGENIC TEMPERATURES'*^

Patrick T. Purtscher

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

Stretch zones are normally observed on the surface of

fracture toughness specimens that break in a ductile manner.

However, no stretch zones were observed in the scanning electron

microscope on the ductile fracture surfaces of austenitic

stainless steels broken in accordance with the ASTM E 813

procedure at cryogenic temperatures.

Key words: stretch zone, fracture toughness testing, ductile

fracture, austenitic stainless steel, cryogenic

properties

.

The purpose of this note is to document an apparent anomaly concerning

the formation of stretch zones. The observation of a stretch zone between

the fatigue-precrack and overload fracture regions of fracture toughness

specimens was first reported by Spitzig^ in 1968. Bates and Clark^ observed

stretch zones on specimens of various carbon and alloy steels, aluminum

alloys, and ferritic stainless steel. They related the size of the stretch

zone to the measured fracture toughness. More recently, the Japan Society of

Mechanical Engineers^ has recommended a multiple-specimen procedure in which

the Jjq value is determined from the size of the stretch zone.

At the National Bureau of Standards in Boulder, a number of elasticn c plastic fracture toughness tests have been conducted on ferritic steels.

Stretch zones have been observed in the scanning electron microscope

(SEM) between the fatigue-precrack and overload fracture regions in all

tests conducted at temperatures above the ductile-to-brittle transition

temperature

.

In addition to the testing of ferritic steels, numerous fracture

toughness tests of nitrogen-strengthened austenitic stainless steels have

been conducted in liquid helium at M K and in liquid nitrogen at 76

These tests were conducted according to the ASTM E 813^ procedure for

^Contribution of NBS; not subject to copyright.

89
determining Jj^ by using the unloading compliance technique to monitor crack

growth. At these cryogenic temperatures , the austenitic steels have yield

strengths between 20Q and 1400 MPa, depending upon the exact composition and

temperature. The mealsured fracture toughness values vary widely, from 80 to

400 MPa/m. Their extremely high toughness at cryogenic temperatures is

attributed to the face-centered cubic (fee) structure of the steel. A ductile-to-brl ttle transition is not normally observed in austenitic steels.

For all the approximately fifty austenitic specimens examined in the SEM,

no stretch zones have been observed.

The reason stretch zones are not observed on the ductile fracture

surfaces of austenitic stainless steels at cryogenic temperatures is not

obvious. Table 1 summarizes the mechanical test results for three samples: a 316LN wrought steel that is completely austenitic, a 308L welded steel that

is predominantly austenitic, and an HY80 welded steel that is ferritic. The

toughness values of all three samples are approximately the same, but their

yield strengths are different. The two examples of austenitic steels were chosen because in one case the yield strength is higher and in the other, the

strength is lower than that of the ferritic steel. In addition, the 308L

transforms to approximately 755^ martensite in the gauge section during a tensile test at 76 K. The 316LN does not transform to martensite in a tensile test at 4 K. Therefore, the absence of a stretch zone is not related

simply to the size of the plastic zone or to the austenite stability.

The fracture surfaces of the three different steels are shown in SEM

micrographs (Figure 1). No stretch zone is visible between the fatigueprecrack and overload fracture regions for the 31 6LN and 308L samples. The

fatigue fracture features appear to be continuous up to the beginning of the

dimpled rupture features characteristic of the overload region. The stretch

zone should be about the same size as the crack-tip-opening displacement

(CTOD) values shown in Table 1 . The HY80 sample shows a distinct region

between the fatigue and overload regions that corresponds to the stretch

zone.

1 n Previous work by Gerberich and Hemmings on metastable austenitic

steels that transform to martensite during plastic deformation showed that

stretch zones were easily observed with an SEM, but were not clearly related

to the CTOD. Schwalbe^ ^ noted the presence of stretch zones in aluminum

alloys and showed that the logarithm of the stretch zone size correlated well

with the logarithm of the stress intensity factor up to the point of stable

tearing.

1 p In tests at room temperature, deVries and Schaap measured stretch zones for ductile crack growth of type-304 stainless steel with a contourograph, but also reported that no distinct separation between blunting and

stable crack growth is visible in the SEM. They considered the crack growth

process in type-304 stainless steel to be "a continuous process of alternating crack-tip stretching and ductile crack extension ... from the very

beginning up to final crack extension." This observation agrees in principle

with our experience in testing stainless steels at cryogenic temperatures.

90
In summary, the observation of a stretch zone is not a necessary

condition for the initiation of ductile fracture. Certainly, in many cases,

stretch zones are observable and may be related to the energy required to

initiate ductile fracture. However, in the case of austenitic steels tested

at cryogenic temperatures, no stretch zones were observed in SEM fractographs .

ACKNOWLEDGMENTS

Most of the testing of austenitic stainless steels at cryogenic

temperatures was supported by the Department of Energy, Office of Fusion

Energy. The 308L material was tested in a program supported by the Mobil Oil

Corporation. The testing of the HY80 steel was supported by the D. W. Taylor

Naval Ship R & D Center.

REFERENCES

1. W. A. Spitzig, A fractographic feature of plane-strain fracture in

0 .45C-Ni-Cr-Mo steel; Trans. ASM, vol. 6l , 1968, pp. 3^^~3^9.

2. R. C. Bates and W. G. Clark, Jr., Fractography and fracture mechanics,

Trans. ASM, vol. 62, 1969, pp. 380-394.

3. H. Kobayashi, K. Hirano, H. Nakamura, and H. Nakazawa, A fractographic

study on evaluation of fracture toughness. Fracture 1977 , Proc. ICF4,

vol. 3, University of Waterloo Press, Waterloo, Canada, 1977, pp. 583-592.

4. F. Ebrahimi, Ductile-to-Brittle Transition in Steel Weldments for Arctic

Structures, NBSIR 85-3020, National Bureau of Standards, Boulder,

Colorado (1985). Also to be published in Proc. of the 1 8th Annual

Fracture Mechanics Conference , American Society for Testing and

Materials, Philadelphia.

5. T. A. Siewert, R. E. Trevisan, and P. T. Purtscher, The effect of

electrode weave procedure on HY-80 GMA welds, submitted to the Weld. J.

for publication.

6. 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--!! , NBSIR 79 - 1609 , National Bureau of Standards, Boulder, Colorado, 1979,

p . 61 .

7 . R. L. Tobler, D. T. Read, and R. P. Reed, Strength/toughness relationship for interstitially strengthened AISI 304 stainless steels at 4 K Temperature, in Fracture Mechanics: Thirteenth Conference , Richard

Roberts, ed., ASTM STP 743, American Society for Testing and Materials,

Philadelphia, 1981, p. 250.

91
8. R. P. Reed and P. T. Purtscher, Nickel and nitrogen alloying effects

on the strength and toughness of austenitic stainless steels at 4 K,

Advances in Cryogenic Engineering- Materials , vol. 32, Plenum, New York,

1986, p. 43.

9. Standard test method for a measure of fracture toughness, ASTM

E 813“81, part 11, American Society for Testing and Materials,

Philadelphia, 1981, p. 763-

10. W. W. Gerberich and P. L. Hammings, Fractographic observations of

stretched regions in front of fatigue cracks, Trans. ASM, vol. 62,

1969, p. 540.

11. K. Schwalbe, Crack propagation in AUZnMgCu 0.5 during static loading,

Eng. Fract. Mech., 1974, vol. 6, p. 415.

12. M. I. deVries and B. Schaap, Experimental observations of ductile crack

growth in type 304 stainless steel, in Elastic-Plastic Fracture Test

Methods; The User’s Experience , ASTM STP 856, E. T. Wessel and F. J.

Loss, Eds., American Society for Testing and Materials, Philadelphia,

1985, p. 183.

92
Table 1 . Mechanical Test Results

alloy

f orm

31 6LN

wrought

308L

weld

HY80

weld

microstructure austenite austenite

+ 3 % ferrite

+^% martensite

ferrite

fracture toughness,

Kj^(J) (MPa/m) 220 21 7 248

yield strength,

Oy (MPa) 1022 244 628

fatigue precracking condition. 33 22 25

at 76 K at 76 K at 295 K

plastic zone size.

Op (mm)** 2.4 to 7.2 42 to 126 8.3 to 25

CTOD (mm)*** 0.12 0.47 0.24

test temperature (K) 4 76 295

stretch zone no no yes

* Not a valid test according to ASTM E-813,^ because the specimen thickness was less than 25 Jg/Oy

** Tp (plane strain) = (1 /3)[Kj^^/(2TTOy^)] and (plane stress) = Kj^^/( 27rOy^)

.

*** CTOD (crack-tip-opening displacement) = ( 1 /2) [K^/( EOy ) ]

.

93
316LN 308L

Overload

Fatigue

Overload

50

Stretch zone

Fatigue

HY80 (c)

Figure 1 . SEM photomicrographs of fracture toughness specimens from

a) 316LN, b) 308L, and c) HY-80. All photos were taken

with a 0“^ tilt angle.

94
X

w^’,-

Xi-

-fi,. f.l %|^ > jr& ,&oTwfr)6tlti'.ii’ »iyt«»^x^' >^,v

tUk» « I**./ ^ V*

*, H!Oui^i<?t*J.'-C^'9ra<2c .,- - ., -g:-: -. .,/•

'>U

-^^ ''

i

'. /

'‘'

•

' ' —--. ^

<& ‘t 4* ‘ip-t '.' '^' V ‘ ; ,y.

' 4 .'yi*-,''- ,'i^

jif - ' # t' 'I'"a ;:-'%rG25i^)Jpoi%^4}3:^l«1i^^ V-;S.-%p«fctir«ntW ' *

' ''>'> ^ '4: '

1^.4^

LOW TEMPERATURE DEFORMATION OF COPPER AND AN AUSTENITIC STAINLESS STEEL*'*'

R. P. Reed and R. P. Walsh

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The tensile-deformation characteristics and effect of strain rate were

studied on relatively pure CDA 102 Cu and solid-solution-strengthened AISI

310 , Tensile strain rate was varied between two orders of magnitude (2 x 10“3

2 X 10”^ s~^ ) at temperatures ranging from 4 to 295 K. Tensile stress-strain

and strain-hardening curves were determined for these temperatures. The effect of strain-rate changes on tensile flow strength was measured from strains

near 0.002 (yield strength) to over O. 3 OO. The data reflect three distinct

ranges of face-centered cubic, polycrystalline plastic deformation, which

have different characteristics depending on solute content.

INTRODUCTION

Superconducting magnet structural materials deform at greatly different

rates, varying from relatively slow stress increases due to rising magnetic

fields to quite high strain rates due to quenching effects. Controlled

strain-rate experiments not only provide data on the effects of strain rate

on tensile properties, but they also lead to a better understanding of deformation mechanisms. Guntner and Reed^ varied the strain rate from 0.0008 s””' to 0.017 s

^ in separate tests to assess the low temperature (4-295 K) tensile properties of various bar, sheet, and plate forms of alloy 304. Variations of yield strength up to 10^ and of tensile strength up to 25^ were observed from increasing the strain rate for this metastable austenitic

stainless steel.

The dependence of the strain to initiate load drops on the strain rate

during tensile testing of alloy 304L was measured by Read and Reed^ at 4 K. As the strain rate was increased from about 0.003 to 0.1 s~\ the strain

prior to the first load drop decreased from about 0.05 to 0.01^. From thermocouples embedded within tensile specimens, temperature rises up to 52 K were detected during load drops. Load drops are associated with macroscopic

slipbands, closely resembling the Liiders bands that appear during ambient

temperature tensile tests. The temperature rise is associated with the very

^Partially supported by the Office of Fusion Energy, U.S. Department of

Energy. Contribution of NBS, not subject to copyright.

To be published in Advances in Cryogenic Engineering—Materials, volume 32.

95
low specific heat and low thermal conductivity of metals at very low tempera

tures, which result in nearly adiabatic localized conditions. Using fine

superconducting wire bonded within a slot on the surface of tensile specimens, Read and Reed did not detect temperature rises up to 10 K (T^ of the

NbTi wire) during plastic deformation prior to load drops. They concluded

that the triggering event for discontinuous yielding in a macroscopic sense

was essentially mechanical in nature and that only during the dislocation

avalanche, associated with the load drops, was the temperature adiabatically

driven up to much higher temperatures. The microscopic triggering event for

the macroscopic slipband (dislocation avalanche) has not been adequately

identified.

Conrad^ reviewed the low temperature deformation of face-centered cubic

metals. He analyzed existing strain-rate-dependent data of Al, Ag, Cu, and

Ni, plotting a strain-rate sensitivity parameter (AJlno/Ailne) versus temperature over the range 4 to 300 K. He illustrated that AS,no/A£ne gradually de- creases at lower temperatures, falling to zero at 0 K for these metals.

Conrad showed that the activation volume (reciprocal of stress sensitivity)

decreases as a function of stress (and strain) at constant temperature and

that the activation volume of Cu is greater than that of Al.

Tien and Yen^ summarized work on low temperature creep. The review includes some of the tensile measurements on Cu that use strain-rate-change

experiments to measure activation volume. The apparent disparity of experimental results on this subject is compounded by the dependency of activation

volume on flow stress (o), strain (e), and temperature (T). There have been

no low temperature strain-rate-change measurements on austenitic stainless

steels; however, Warren and Reed^ published tensile stress-strain curves at

low temperatures on alloy 310.

Kocks et al.^ ^ and Mulford”'^ presented a clearer representation of

strain hardening, strain-rate sensitivity, and dynamic aging for polycrystalline metals and alloys. They argued that the flow strength at low temperatures

is composed of two additive contr ibutions--from dislocation-solute interactions and dislocation-dislocation interactions. Their approach is discussed

in more detail later and used to describe some of our data.

This paper presents stress-strain , strain-hardening, and strain-rate

sensitivity data at temperatures from 4 to 295 K for relatively pure Cu and

for a solid-solution strengthened austenitic stainless steel.

^EXPERIMENTAL PROCEDURE

Specimens of relatively pure Cu (oxygen-free, high-conductivity CDA 102)

and a common solid-solution-strengthened austenitic stainless steel (AISI

310) were obtained in 1 .9-cm (0.75-in) diameter bar stock and tested in the

annealed condition. Alloy 310 was annealed at 2050°C for 0.5 h after machining and CDA 102 Cu, at 650°C for 1 h. Round tensile specimens were machined

to a 6 . 35 -mm diameter along a gage length of 41.9 mm, increasing with a radius of 9.52 mm to 12.7-mm diameter threaded ends. The entire specimen

length was 71 .4 mm. The hardness, grain size, and chemical composition of

both alloys are listed in Table 1 .

96
Two loading frames were used. At room temperature the loading compliance of the screw-driven assembly was measured as 8 x 10~3 cm/kN (8 x

10”'^

cm/kgf). The compliance of the servohydraulic load train varied from 11 x 10'^ cm/kN at loads less than about 2 kN to 6 x 10“3 cm/kN at larger loads.

The Cu specimens were measured at the lower compliance levels of the load

train; the austenitic stainless steels, at the higher compliance levels.

The actuator velocity and crosshead movement of the machines were independently measured and found to deviate from expected rates. The "low"

crosshead rate of the screw-driven machine was 0.0047 cm/min (stated as 0.005

cm/min) and the "high" crosshead rate for this machine was 0.455 cm/min

(stated as 0.5 cm/min). The servohydraulic "high" actuator speed was 0.540

cm/min (stated as 0.5 cm/min), significantly higher than that of the other

machine. The "low" actuator speed of the servohydraulic machine was 0.005

cm/min .

In all tests, specimens were initially loaded at the low strain rate

(2 X 10"^ ) at constant temperature. After sufficient strain to permit

measurement of the yield strength (0.002 plastic strain), the strain rate (e) was incrementally increased (to 2 x 10~^ s~^ ) and held long enough for good

character ization of the stress-strain curve. The strain rate was subsequently reduced, and the specimen was strained to the next measurement range.

The difference between the flow strengths at equivalent strain-hardening

rates was considered the stress change (Ao) associated with the strain-rate

change. The exception to this was measurement of alloy 310 at higher levels

of strain, where apparent "yield points" were observed immediately following

changes of strain rate. In these cases we used the maximum value of Ao. Use

of extrapolated Ao to the onset of the high strain-rate range usually produced a negative dependence (but always positive values) of Ao on e or o.

Temperature measurement and control, measurement equipment variables

that affected the accuracy of Ao data from strain-r^te changes, strain mea- surement, and other procedures are discussed in Ref. 12. Reference 12 also

contains all data from which the strain-rate sensitivity figures of this

paper were derived.

At low strains (-<0.04) the attainment of an equivalent strain-hardening

rate on changing to the higher strain rate usually consumes more strain.

Therefore, interpretation of Ao at lower o or e values is more difficult.

The strain range over which the strain-hardening rates tended to be slower to

change was always less than 0.01. This early portion of the o-e curve naturally exhibited larger strain-hardening rates and, when reliable measurements

could be obtained, lower values of Ao. No data in this low strain range were included unless both o-e curves (high e and low e) assumed equivalent strainhardening rates.

Discontinuous yielding at temperatures below about 20 K was observed in

both materials. The strain at which load drops began was much less in the

austenitic steel than in Cu. This is atcributed to their differences of low

temperature thermal conductivity. The thermal conductivity of alloy 310 at

4 K is about 0.5 W-m"'' •K'"' , and that of Cu is 66? W*m“’' •K"'' . The specific

97
heat of the austenitic steel is actually somewhat higher (~ 1.8

than that of Cu (0.10 J*kg~^*K“^). The occurrence of discontinuous yielding

precluded assessment of strain-rate effects and strain-hardening rates on flow stress. Changes to the higher strain rate immediately triggered load

drops and more erratic stress-strain behavior. At the higher strain rate it

is possible that the flow strengths at 4 K reflect thermal instability, even

at lower strains prior to the onset of discontinuous yielding.

RESULTS AND DISCUSSION

Copper

The tensile stress-strain curves of CDA 102 Cu are shown in Fig. 1. At

low temperatures there was a decided increase in strain hardening but little

increase in yield strength. The tensile yield strength Increased from about

27 MPa at room temperature to 32 MPa at 4 K. Only very modest discontinuous

yielding occurred at 4 K at large strains.

The data depicting the dependence of Ao on o for Cu are summarized in

Fig. 2. At 4 K there is a linear dependence of Ao on a. At higher temperatures the dependence of Ao on o is linear at low strain and slightly concave

upwards at higher strain. The Ao decreases with decreasing temperature and

extrapolates to zero at zero stress.

Alloy 310

Alloy 310, a stable austenitic stainless steel, has completely different

stress-strain and strain-rate sensitivity traits. After deformation neither

martensitic product (hexagonal close-packed e; body-centered cubic a’) was detected at any temperature. The stress-strain curves from 4 to 295 K are shown in Fig. 3* Alloy 310 shows considerably more strength in the early re- gion of the stress-strain curve than relatively pure Cu Indeed, the tensile

yield strength increased from about 215 MPa at 295 K to 730 MPa at 4 K. The

rate of strain-hardening of alloy 310 at high strain was rather independent

of temperature, unlike that of Cu.

Data on Ao/AHne versus a for alloy 310 are summarized in Fig. 4. Notice

the differences first between Cu (Fig. 2) and alloy 310. The Ao values of

alloy 310 were nearly constant, independent of stress or strain and, extrapolated to 0 = 0, always are positive. The maximum of Ao as a function of temperature occurred in the range 76 to 100 K. From 76 to 4 K, there was a sharp decrease of strain-rate sensitivity.

At strains in excess of about 0.02 at temperatures near 76 K and strains

of about 0.05 at temperatures near 295 K, very small yield points began to occur in alloy 310 following strain-rate changes. Again, the maximum Ao values on changing from a low to high strain rate were plotted. Extrapolation

of equivalent strain-hardening rate ranges to the strain at the time of

strain-rate change yielded values up to 15^ less than the reported data. Use

of these data of lower Ao would have resulted in slightly negative slopes of

Ao versus o or e.

98
The Ao data at 4 K exhibit a negative slope; however, as discussed earlier, these data may reflect partial thermal instability at the higher strain

rates. After changes from low to high strain rate, a brief region (~0.0003e)

of increased flow strength was observed, followed by a load drop. Although

the maximum Aa was recorded for each strain-rate change, these Ao data should

be regarded as minimal; for it is possible that internal specimen heating at

the high strain rate effectively reduced the maximum flow strength.

Stress-Strain Characteristics

^“11 Following the general approach of Kocks,° the flow strength (a) can

be expressed as:

0 = 0^ + 0^ ( 1 )

where is the frictional strength contribution dependent upon dislocationsolute interactions, and is the dislocation-dislocation interaction con- tribution dependent primarily on number and structure of dislocations. The

frictional strength can be regarded as independent of strain; the yield

strength is a good measure. The term is strain dependent; tensile strainhardening characteristics are reflective of its contribution. Typically,

Ojj = ayb/p (2)

where u is the shear modulus; b, the Burgers vector; p, the active dislocation density; and a, a constant, which may be dependent on the solute content, c, and is slightly less than unity. In face-centered cubic alloys, the

strain-hardening rate (do/de) is dependent on solute content. To account for

this. Bloom et al.^ suggest

0 ^

= [1+k(c)] ayb/p (3)

with k(c) expressing the effect of solute. The strain-hardening rate is

expressed as

0 = ^ ^ - 1/2 ^

de 2 ^ de ^ ^

The change of dislocation density with strain is considered to have two components: one concerns the "storage rate" and the other, the dynamic recovery

rate. The rate of dislocation accumulation is athermal in nature; the recovery rate is obviously affected by thermal activation. Thus, with these con- siderations, the strain-hardening rate should have two major terms:

0 = da/de = 0j.j(c) - 0j^(e,T,o,c) (5)

where 0j^ is the athermal hardening rate and 0^ is the dynamic recovery rate

contributing negatively to 0.

The strain-hardening curves derived from the stress-strain curves for

Cu are shown in Fig. 5. They are plotted as a da/de versus o, following

99
Mecking and Kocks.'^ There is a linear region with positive slope at low

stress followed by deviation to lower strain-hardening rates at higher

stress. The linear region is constant at- all temperatures and deviations

from linearity occur at higher stresses with lower temperatures.

The strain-hardening curves (Fig. 6) at various low temperatures of alloy 310 indicate two distinct stages. The strain at the transition of the

two stages increased from about 0.02 at 295 K to 0.04 at 4 K. This is,

approximately, the restricted strain range that separates the early (difficult-to-measure ) stage of strain-rate sensitivity from the higher strain re- gion. The strain-hardening in the early strain range was very sensitive to

temperature, similar to the yield strength. The strain-hardening rates in

the higher strain range were relatively insensitive to temperature and

assumed the same slope, which was negative, but only slightly. In fact, if

the curves were normalized by dividing the ordinate by there would be even less difference in their absolute magnitude. There was no evidence of

dynamic recovery, since a recovery stage should be sensitive to temperature

and nonlinear.

Thus, if Cu and alloy 310 are representative of pure and solid-solutionstrengthened face-centered cubic metals, one may make the following conclusions about their strain-hardening characteristics at low temperatures: There

is an early stage of rapidly decreasing high rates of strain-hardening in the

region of the yield stress for solid-solution-strengthened alloys; this re- gion is strongly temperature dependent. The next stage (first stage for pure

metals) exhibits a linear dependence of o0 on o and, for Cu, is athermal. The

slopes of this stage depend on solute content (solute effects on dislocation

interactions); in pure metals the strain-hardening dependence extrapolates to zero (indicating only dislocation-dislocation interactions). The final stage

is a region of dynamic recovery, not necessarily present in solid-solutionstrengthened alloys. This stage typically exhibits a nonlinear dependence of

00 on 0 .

Strain-Rate Effects

Alloy 310, with much larger flow strength than Cu, exhibited strength

changes of the order of 20 to 40 MPa with each change of strain rate; Cu had

strength changes of less than 5 MPa. In both alloys, the stress change approached zero at very low temperatures. However, if the data were normalized

by dividing the stress change by the yield strength at temperature, all stress

changes would be found to be less than 155? of the flow strength. All percentages decreased to zero at low temperatures. There is a distinction between measurements of strain-rate effects from experiments that maintain a given strain rate and those that change strain rate during straining. During

continuous straining, higher strain rates induce internal specimen heating

more rapidly. This heating cannot be quickly removed at the high strain rate;

thus the specimen probably remains permanently at a higher temperature during

the test. This results in lower flow strengths and failure strengths. On

the other hand, the strain-rate change experiments provide time at the low

strain rate for the specimen to return to thermal equilibrium and, thus, to

return to the higher flow strength at 4 K.

100
Mulford and Kocks^*^ have shown that, provided the sum of the contributions of solution strengthening (o^) and work hardening (a^) are equivalent

to the flow strength, the strain-rate sensitivity has two components;

3o

3ilne

Of (m^ - m^) + 0 m^j ( 6 )

where m^ = Dfi.n0f'/9ilne and m^ = D£no^/3£ne at constant temperature. A plot of

9o/3Jlne versus o, called the Haasen plot, should have a slope of m^ at stresses somewhat above and extrapolate to o^(m^ - m^^) at zero stress. If two

contributions are present, they should be reflected in a change of slope at

rather low stress: at low stress the contribution from dislocations must be

small anc^, therefore, the contribution from solute strengthening dominates.

The Cu data of Fig. 2 extrapolate to zero as flow stress approaches zero. This indicates the dominance of dislocation-dislocation interactions and the

absence of solute strengthening (o^ = 0). For alloy 310, the range of con- stant Ao/AZne values corresponds to the second stage of oO versus a (Fig. 3),

where o0 is relatively insensitive to a.

Alloy 310 exhibited considerable solute strengthening, and extrapolation

of the Aa/AZne data to zero stress yielded large values of frictional stress

(oj> is positive), dependent on temperature. Unlike Cu, the slope of Ao/AHne

was practically independent of flow stress. Therefore, the solute-strengthening mechanism operative in this second stage (related to strain-hardening

behavior) is very rate sensitive and temperature dependent. Both strainhardening and strain-rate data confirm the absence of dynamic recovery in

alloy 310 at the measured o levels. There is a distinction between the ratelimiting process that controls the strain-rate response at low strain levels

and the process that influences the response at higher strain, and it appears

that both are influenced strongly by solutes. At low strains, perhaps the

interaction of individual dislocations with discrete solutes is rate controlling and short range. At high strains, perhaps dislocation cell wall interaction with solutes or dislocation multiple solute interactions are rate

controlling.

The temperature dependences of both strain-hardening and strain-rate

sensitivity showed no unexpected transitions; therefore it appears that the same rate-controlling deformation mechanism that operates at room temperature

remains operative at very low temperatures in both metals.

SUMMARY

The stress-strain curves, strain-hardening curves, and strain-rate sen- sitivities have been determined for Cu and alloy 310 austenitic stainless

steel at temperatures from ^ to 295 K. The results suggest three regions of

plastic deformation of polycrystalline face-centered cubic metals and alloys:

101
1. For solute-strengthened alloys at low strain, a region of rapidly

decreasing strain-hardening rate and rapidly rising strain-rate sen- sitivity is strongly dependent on temperature and solute content.

2. At low strain for pure metals and at intermediate strains for solutestrengthened alloys, a strain-hardening region where the slope of a0 versus a is constant and strongly dependent on solute content. In

this region, the strain-rate sensitivity plotted against a has a positive or zero slope, dependent also on solute content.

3. At high strain, a region of dynamic recovery in metals and some alloys. Here, the a0-versus-o trend shows an ever-increasing negative

slope, and the plot of strain-rate sensitivity versus stress shows

an ever-increasing positive slope.

Finally, there is less strain-rate sensitivity at very low temperatures,

although specimen adiabatic heating at higher strain rates does result in the

lowering of the flow stress.

REFERENCES

1 . C.

2. D.

3. H.

4. J.

5. K.

6. U.

7. H.

8. R.

9. T.

10. R.

11. R.

12. R.

J. Guntner and R. P. Reed, ASM Trans. Q . 55:399-^19 (1962).

T. Read and R. P. Reed, Adv. Cryo. Eng. —Mater . 26:91-101 (1980).

Conrad, in "High Strength Materials," J. Wiley & Sons, New York

(1965), pp. 436-505.

K. Tien and C.-T. Yen, Adv. Cryo. Eng.—Mater . 30:319-338 (1984).

A. Warren and R. P. Reed, "Tensile and Impact Properties of Selected

Materials from 20 to 300 K." NBS Monograph 63 » U.S. Government

Printing Office, Washington, D.C. (1963).

F. Kocks, R. E. Cook, and R. A. Mulford, Acta Metall. 33:623-638

(1985).

Mecking and U. F. Kocks, Acta Metall . 29:1865-1875 (1981).

A. Mulford and U. F. Kocks, Acta Metall . 27:1125-1134 (1979).

A. Bloom, u. F. Kocks, and P. Nash, Acta Metall . 33:265-272 (1985).

A. Mulford and U. F. Kocks, Scripts Metall. 13:729-732 (1979).

A. Mulford, Acta Metall . 27:1115-1124 (1979).

P. Reed and R. P. Walsh, Low Temperature Tensile Characteristics of

Copper and Aluminum, in "Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures —IX," NBSIR 86-3050, National Bureau

of Standards, Boulder, Colorado (1986), pp. 107*140.

102
Table 1 . Material Characterization

Alloy Chemical

Composition

(wt.^)

Hardness Grain

Size

(pm)

CDA 102 99.99 Cu + Ag 22 36

310 211.8 Cr, 20.8 Ni, 1.7 Mn, 0.093 C,

0.031 N, 0.02 P, 0.02 S, 0.7 Si,

0.1 Mo, 0.1 Cu, balance Fe

72 1 61

Fig. 1. The tensile stress-strain curves of CDA 102 Cu at various

temperatures.

103
STRESS,

MPa

Fig. 2. General character ization for CDA 102 Cu of the effects of strainrate change (2 x 10 “^ to 2 x 10 “^ ) on the stress change as a function of flow stress at various test temperatures.

Fig. 3- The tensile stress-strain curves of alloy 310 at various

temperatures.

104
10

Fig.

Fig. 5.

Alloy 310

cf CL 6 —

CO

CO

LJ

q:(-CO<

CM

•'V c

Mlc

/295 K

76 K

.50 K

2 —

^25 K ^20 K

•4 K

400 800

STRESS, MPa

1200

General characterization for alloy 310 of the effects of strainrate change (2 x 10“^ to 2 x 10~3 ) on the stress change as a function of stress at various test temperatures.

Strain-hardening curves for CDA 102 Cu derived from stress-strain

curves of Fig. 1.

105
Fig. 6. Strain-hardening curves for alloy 310 derived from stress-strain

curves of Fig. 3 .

106
u;- , ^ 'SKSiiSl

lU

%

*

V;, .' V'. ^'*". ' r^W/'

Sf*^. i

5ft"

•-»r •j^^/diifn- t]rt<s 'r*i,rili: Sas:-'. ' ’3 *1

• r ->1RL« in. -i-i-,. _. ...,. . -‘^i^"'^-

rp,:"V;Kt I4i .,. ..•> .f

LOW TEMPERATURE TENSILE CHARACTERISTICS OF COPPER AND ALUMINUM

R. P. Reed and R. P. Walsh

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The tensile-deformation characteristics and effect of strain rate were studied on relatively pure copper and aluminum alloys, CDA 101 Cu and

1100 A1 . Tensile strain rate was varied by two orders of magnitude

(2 X 10”^; 2 X 10~5 s”"* ) at temperatures ranging from 4 to 295 K. Tensile

stress-strain and strain-hardening curves were determined for these temperatures. The effect of strain-rate changes on tensile flow strength was measured from strains near 0.002 (yield strength) to over 0.300. The data

reflect distinct ranges of face-centered cubic, polycrystalline plastic deformation; their particular characteristics depend on solute content.

INTRODUCTION

Superconducting magnet structural materials deform at greatly different

rates, varying from relatively slow stress increases due to rising magnetic

fields to high strain rates due to magnetic-field quenching effects. Controlled strain-rate experiments not only provide data on the effects of

strain rate on tensile properties, but they also lead to a better understanding of deformation mechanisms.

Conrad”' reviewed the low temperature deformation of face-centered cubic

metals. He analyzed existing strain-rate-^dependent data for Al, Ag, Cu, and

Ni by plotting a strain-rate sensitivity parameter ( d2,na/d£ne ).p versus temperature, T, over the range 4 to 300 K. He illustrated that d2,no/dS,ne gradually decreases at lower temperatures, falling toward zero at 0 K for these

metals. Conrad showed that the activation volume (reciprocal of strain-rate

sensitivity) decreases as a function of stress and strain at constant temperature and that the activation volume of Cu is greater than that of Al

.

The activation volume is expressed as the reciprocal of strain-rate sensitivity: V = kT(dfi,ne/da)Y. Warren and Reed^ published tensile stress-strain

curves at low temperatures on CDA 101 Cu and 1100 Al.

Kocks et al.^ and Mulford^ presented a clear representation of strain

hardening (do/de), strain-rate sensitivity ( do/d£ne ),p , and dynamic aging for

polycrystalline metals and alloys, where o is tensile stress and e is tensile

strain. They argued that the flow strength at low temperatures is composed

^Partially supported by the Office of Fusion Energy, U.S. Department of

Energy. Contribution of NBS, not subject to copyright.

107
of contributions from dislocation-solute interactions and dislocationdislocation interactions. Their approach is discussed in more detail later

and used to describe some of our data.

The strain-rate sensitivities of the tensile deformation of copper have

been measured for single crystals^'^^ and polycrystals . ^ ^ Butt and

Feltham^^ and den Otter and van den Beukel”'^ studied polycrystalline behavior at temperatures as low as 77 K. Chr istodoulou and Jonas”' ^ and Regazzoni

and Montheillet^ reported the strain-hardening characteristics of their

polycrystalline copper. Mecking and Kocks studied silver.^ Butt and Feltham

found an inverse linear relationship between the log of the activation volume

and applied stress between 77 and 290 K. Regazzoni and Montheillet observed

a linear dependence of the strain rate sensitivity on applied stress (6O-3OO

MPa range) at room temperature and reported that do/diine extrapolates to zero

at zero stress. A similar trend was observed by Melander and Thuvander^^ in

the stress range 200 to 330 MPa. However, Chr istodoulou and Jonas reported a minimum in the strain-rate sensitivity versus applied stress at room temperature located between 230 MPa for large strain-rate changes (e = 5 x 10~^ to

10

-2 ^-1

s ') and 410 MPa for smaller changes.

There have been a number of studies to measure the strain-rate sensitivity of A1 in single”'^ and polycrystals . ® ^ ^ Some values were measured at temperatures as low as 4 others, to 77 k.

Saimoto et al.^^ reported a minimum of the term (1/T)[Ao/(o - Oy)/(jinGi - S,ne2)] versus temperature at about 225 K from their strain-rate change mea- surements at temperatures from about 78 to 400 K, where 0 is the yield

strength at 0.002 strain offset. Basinski^^ reported a plateau of

( 1 /T ) ( do/dJlnl ),p over the temperature range 10 to 80 K, followed by a de- crease from 80 to 295 K.

Saimoto et al 22-26 and Mulford and Kocks5,7,1 also considered the ten-

,22 slle strain-hardening characteristics. Saimoto, Sang, and Morrls^"^ found a convergence of the strain-hardening term (a - ay)(Aa/Ae) as the adjusted flow

stress a - a ) decreased and an increasing range of linearity of this relationship as the temperature decreased. This dependence is similar to that

reported by Mecking and Kocks

polycrystalline Cu.

for Ag and by Bullen and Hutchinson^^ for

This paper presents stress-strain , strain-hardening , and strain-rate

sensitivity data at temperatures from 4 to 295 K for a relatively pure

(99.99^) Cu and for a relatively pure (99.1^) solid-solution strengthened Al.

EXPERIMENTAL PROCEDURE

The hardness, grain size, and chemical composition of the Cu and Al are listed in Table 1 . Both materials were obtained in 1 .9-cm bar stock and

tested in the annealed condition. The CDA 101 Cu (oxygen-free, high conductivity) was annealed at 650°C for 1 h in Ar; 1100 Al, 500°C for 0.5 h in air;

both were air cooled. The residual resistivity ratios were 105 for the CDA

101 Cu and 35 for the 1100 Al alloy following annealing. Round tensile

specimens were machined to a 6.35-mm diameter along a gage length of 41 .9 mm,

tapering with a radius of 9.52 mm from 12.7-mm threaded ends. The length of

the entire specimen was 71.4 mm.

108
The cryostat equipment is shown in Fig. 1 . Tests at 4 and 76 K were

conducted in boiling liquid helium and nitrogen, respectively. Other low

temperatures were achieved by testing in helium (^-76 K) or nitrogen (>76 K)

gas and automatically adjusting the gas flow and the heater input at the

specimens grips with temperature controllers.

The type E thermocouples were taped to the middle of the specimen on a varnished surface and insulated from the gas atmosphere with foam. Silicon

diodes inserted into small drilled holes in both specimen grips permitted

digital readout and control of temperatures at both the upper and lower

grips. By these means, we were able to maintain specimen temperature within

1.0 K during testing at 20 K (other temperatures were easier to control).

During discontinuous load drops while testing at temperatures from 4 to

20 K, surface temperature increases up to 50 K were detected. Specimen temperature returned to the desired control temperature during subsequent elastic deformation to the prior flow-stress level.

In all tests, the specimens were initially loaded at the low strain rate

(2 X 10'^ s~”’ ) at constant temperature. After sufficient strain to permit

measurement of the yield strength, the strain rate was abruptly increased to

2 X 10~3 s~^ and held long enough for good characterization of the stressstrain curve. The strain rate was then reduced, and the specimen was strained to the next measurement range. Each strain increment was about

0.002; this was sufficient to determine the yield point or unloading effects.

However, at temperatures less than about 50 K, specimen warming was observed

during the high strain-rate segments. This warming became detectable at

strains in excess of about 0.001; therefore, smaller strain increments were

used at these low temperatures.?

New, high elongation strain-gage extensometers that have been developed

for use at low temperatures were used for most of these measurements. These

extensometers, capable of sensitivities of 10 have three flexing A1 beams.

The geometry of the beams permits sensing of specimen strain to elongations

of over 3055 at low temperatures. Four 350-P strain gages were mounted to

each beam and wired in a Wheatstone bridge configuration. A 5-V dc power

supply was used to adjust excitation.

Two types of tensile machines were used. With the Instron Model TTDML,*

the cryostat assembly (Fig. 1) uses the downward motion of a screw-driven

crosshead to impart load through a stainless steel compression cylinder to

the specimen. A servohydraulic MTS Model 8 IO,* with the actuator repositioned

above the crosshead to accommodate a more efficient cryostat system, was also

used, but only at temperatures of U, 76 , and 295 K. In this assembly two

fiberglass-epoxy compression posts sustained the load imparted through the

actuator to the specimen. This low temperature loading system is more extensively discussed by Read and Tobler.^”'

*Trade names are used here to describe specifically the method

including loading, compliance, and portions of the load train;

NBS endorsement of these products.

of testing,

this is not an

109
At room temperature, the loading compliance of the screw-driven assembly

was measured as 8 x 10~3 cm/kN. The compliance of the servohydraulic load

train varied from 11 x 10~^ cm/kN at loads less than about 2 kN to 6 x 10 “^ cm/kN at larger loads. The Cu and A1 specimens were measured at the higher

compliance levels of the load trains.

The velocity of the actuator and crosshead movement of the machines were measured independently and found to deviate from expected rates. The "low"

crosshead rate of the screw-driven machine was 0.00^7 cm/min (stated as 0.005

cm/min) and the "high" crosshead rate for this machine was 0.455 cm/min

(stated as 0.5 cm/min). The "high" actuator speed of the servohydraulic

machine was 0.540 cm/min (stated as 0.5 cm/min); its "low" actuator speed was

0.005 cm/min.

The responses of the crosshead, actuator, and load and strain recordings

to change of strain rate were also assessed. Measured with a digital oscilloscope, the transition with respect to crosshead and actuator stability from

low strain rate to high strain rate was quick; the fast strain rate was obtained within the 10-ms sensitivity. However, the transition in the screw- driven machine from fast to slow was much less rapid, and a pause in

crosshead motion was observed for time periods from 0.1 to 2 s. This slackening of crosshead motion usually resulted in a dip in stress (indicating

apparent strain softening) at the onset of the lower strain rate (illustrated

in Fig. 2a) . Therefore, all recorded values of Ao were measured from increasing strain-rate changes. In both machines, the response time of the

strain amplification circuits was adequate to distinguish strain changes from

rate changes.

The time constant (cutoff frequency) for the filtering of the load signal of the two machines differed. The servohydraulic machine had a shorter

time constant (higher cutoff frequency), and its load response to strain-rate

changes of two orders of magnitude was always adequate. The screw-driven

machine had a longer time constant, and its load response sometimes lagged

the strain response during strain-rate changes. This led to possible ambiguities in Interpreting the load-strain data (see Fig. 2b). Particularly, if

the rate of strain hardening was large, the load lag during the rate change

produced apparently low or even negative stress-change values. High strainhardening rates are predominant in Cu and A1 at low strains (<0.04); therefore, data from the screw-driven machine at low strains are not Included in

this paper.

In the subsequent presentation of data, Ao is plotted versus a for Cu

and A1 . Here the flow stress, o, is taken as the average stress of the two

strain rates. Copper and aluminum data from the screw-driven machine were adjusted to match the data at 4, 76, and 295 K obtained from the servohydraulic machine by considering their different strain rates, compliance, and

other machine variables.

Following changes of strain rate, yield point effects were observed at

temperatures ranging from 76 K to room temperature. At lower temperatures

transient effects were not detected until there was considerable strain

(-0.10). Near room temperature transient yielding effects were observed

110
almost immediately. In Figure 3a and b, typical strain-rate change sequences

depict the observations. The difference between Ao and Aa^ increased with

increasing strain; at the end of the test, Ao^ was approximately 105^ of Ao

for A1 and 20^ of Ao for Cu.

We chose to report Ao from our measurements. We obtained Ao by measuring the maximum load difference (the difference between the peak load and the

extrapolated load at lower strain rate; see Fig. 3a) and dividing by the

original specimen area. The value, Ao, was chosen primarily for the following reason: during straining at the higher strain rate at temperatures at

and below 76 K, specimen warming was detected by our temperature monitors.

This made distinctions between transient yield-point behavior and warming

effects difficult to sort out. Specimen warming was never detected in the

early yield-point region.

When changing to the higher strain rate at low strains (-<0.00^0, an equivalent strain-hardening rate was usually not attained. Therefore, interpretation of Ao at equivalent strain-hardening rates was not possible.

Extrapolation to the strain at the point of strain-rate change yielded lower

values of Ao (Fig. 3c). It is likely that in this early strain range the

strain rate influences the rate of dislocation accumulation per unit strain.

No data in this low strain range were included unless both o-e curves (high e and low e) assumed equivalent strain-hardening rates.

Discontinuous yielding at ^ K was observed in Cu at high strains (see

Fig. 3d). Thermal conductivity of the A1 was about 22 W*m~'*K~^ and that of

Cu was 667 W*m*^ *K~”' . The specific heat value of Cu (0.10 J*kg~^*K~^) was lower than that of A1 (0.28 J*kg~^‘K~^). Discontinuous yielding occurred at

about 350 MPa at a strain of 0.32; at these stress and strain levels, 1100 A1

had smooth stress-strain behavior. Therefore, the occurrence of discontinuous yielding in Cu must be closely associated with the lower specific heat

contribution to adiabatic heating and not critically affected by the difference in thermal conductivity. Changes to the higher strain rate immediately

triggered load drops and more erratic stress-strain behavior (Fig. 3©).

Changes from high to low rate resulted in load increases until the steady

state of the low rate was reached (Fig. 3©)* The occurrence of discontinuous

yielding precluded assessment of strain-rate effects on flow stress.

EXPERIMENTAL RESULTS

Copper

The tensile stress-strain curves of Cu are shown in Fig. M. At low

temperatures there was a decided increase in strain hardening but little in- crease in yield strength. Only very moderate discontinuous yielding

occurred at 4 K at large strains. The temperature dependence of the yield

strength followed the same form as face-centered cubic alloys the log of

the strength decreased with increasing temperature:

Oy = 36 exp (- 0.28 X 10~3 Y) . ( 1 )

Therefore, there was a very small dependence of 0 on T from Oy = 36 MPa at

4 K to 33 MPa at 295 K.

y y

111
The data depicting the dependence of Ao on o for Cu from strain-rate

changes are presented in Figs. 5 through 7 for various low temperatures

.

Data in Fig. 5 represent three independent measurements of separate specimens. At 4 K (Fig. 7), there was an apparently linear, very gradual, dependence of Ao on a; at higher temperatures the dependence of Ao on o was slightly concave upward. The stress change. decreased with a decrease in

temperature and apparently extrapolates to zero at zero stress.

The dependence of Ao on e was slightly concave downward, contrasting to

the dependence of Ao on o. The Ao extrapolated to positive values, ranging

from about 0.3 to 1.5 MPa, at zero strain.

Aluminum

The tensile o-e curves of A1 at temperatures from 4 to 295 K are shown

in Fig. 8. As temperature decreased, the rate of strain hardening increased.

There was very little increase of flow strength at low strain levels as temperature decreased. Values of o increased from approximately 25 MPa at

295 K to ^0 MPa at 4 K, following the expression

Oy = 40 exp (-1.67 x 10“3 t) . (2)

The dependence of the Ao resulting from strain-rate changes on o is

shown in Figs. 9 and 10 for A1 at temperatures from 4 to 252 K.

At room temperature the effects of strain-rate change on Ao are not

clear and are not included in this report. Strain-aging significantly influenced the stress-strain characteristics at the low strain rate, since the o-e curve is composed of many small serrations. At the high strain rate, the o-e curve is irregular and the strain-hardening rates are variable. The best

description is that, following an increase of strain rate, Ao shows a very

small increase initially, but during straining, a negative strain-hardening

rate results. This results in a higher Ao if the o-e curve is extrapolated

to the point of strain-rate change or in a lower Ao if some intermediate

strain is used to represent the data.

The dependence of Ao on o increases with increasing test temperature

from 4 to 100 K, and then decreases at higher temperatures. At temperatures

below 100 K, all curves (at large strains) extrapolate to Ao = 0 at zero flow

stress. Above 100 K all curves extrapolate to positive values of Ao at zero flow stress; the extrapolated value Increases with increasing temperature. At

very low temperatures (20 and 4 K), the dependence of Ao on o is linear; at

higher temperatures (2 50 K), the dependence is concave upward at high flow

stress and linear at low flow stress. At very low stress levels (near the

yield strength), measurements are more difficult (owing to lower absolute

stress change) and arbitrary (owing to changing strain-hardening rates with

strain rate). At low stress levels there is some evidence for a second stage

that exhibits greater dependence on flow stress. The data points at the

lowest flow-stress levels of the 20- , 50- , and 252-K plots may reflect the

change of dependency at low stresses.

112
DISCUSSION

Strain Hardening

Following the general approach of Kocks,^"^ the flow strength can be

subdivided:

0 = 0 ^+ , ( 3 )

where is the frictional strength contribution dependent upon dislocationsolute interaction, and is the dislocation-dislocation interaction contribution, dependent primarily on the number and structure of dislocations. The

frictional strength can be regarded as independent of strain, and the yield

strength is a good measure. The term is strain dependent, and tensile

strain-hardening characteristics reflect its contribution. The dislocation

interaction typically is related to the dislocation density.

^d

= ayb p 1/2

(^)

with M, the shear modulus; b, the Burgers vector; p, the active dislocation

density; and a, a constant. The constant, a, may be dependent on the solute

content, c, and is slightly less than unity. Thus, the strain-hardening

rate, 0, is

^ da apb -1/2 dp

dl- — P Tc

and

do

de

= (

0

-

. da

I de

(apb)'

de

(5)

( 6 )

is proportional to the change of dislocation density with strain (dp/de). The

change of dislocation density as a function of strain is considered to have

two components: one is the rate of storage and the other, the rate of dynamic recovery. The rate of dislocation accumulation is athermal in nature;

the recovery rate is obviously affected by thermal activation. With these

considerations, the quantity proportional to the dependence of dislocation

density on strain (Eq. 6), should have two major terms:

a^0 = a^0j^(c) - a^0j^(e,T,a,c) , (7)

where O(j0j^ is related to the athermal hardening rate and o^0p is related to

the dynamic recovery rate contributing negatively to ajj0.

The strain-hardening curves derived from the a-e curves for Cu are

included in Fig. 11. They were derived from Eq. 6 and are plotted as

(a - Oy)Ao/Ae versus a - Oy, with the assumption that a^- = Oy. There is a linear region with positive slope at low stress followed by deviation to

lower strain-hardening rates at higher stress. The linear region is constant

at all temperatures and deviations from linearity occur at higher stresses

with lower temperatures.

113
The strain-hardening data for A1 are shown in Fig. 12. Similar to those

for Cu, there is a linear region of increasing (o - 0)0 with increasing

0 - 0 whose slope is independent of temperature. The slope of this linear

region is slightly greater for A1 (2.2 GPa) than for Cu (1.7 GPa).

The 99.1 A1 is expected to exhibit stronger solute interactions than the

99.99 Cu. These solute interactions resulted in a steeper temperature dependence of Oy for A1 . "Normalizing” by subtracting Oy from both axes produces almost identical sets of curves for Cu and A1 . Saimoto et al.^^ also

used this procedure; Kocks et al.'^”'^ suggested dividing by the shear modulus.

This latter procedure would result in discrete, independent curves as a function of temperature for Al, not in convergence to an "athermal" dependence.

Thus, the derivation and use of Eq. 6 to describe strain-hardening behavior

is reinforced by the Al data.

Dynamic recovery, as evidenced by the deviations from linearity at

higher 0 values, was apparent for both metals. In this region, the rate of

accumulation of active dislocations is decreasing, but always remains positive.

Strain-Rate Effects

When the two major contributions to the flow strength from solute effects and dislocation interactions are taken into consideration, the strainrate sensitivity has two components:

3 g

3S,ne

where m^

Of(m^ - m^j) + om. ( 8 )

3S-na , d

9!lne

at constant temperature.

A plot of 3a/9£ne versus a has a slope of m^^ at stresses above Oy and

extrapolates to o^(m^ ~

'^d^ zero stress. When both contributions are present, they may be distinguished in the low stress range: at low stress

the contribution from dislocation interactions ( 0 ^) is small, and therefore,

the contribution from solute strengthening ( 0 ^) dominates.

The strain-rate sensitivity data of Cu (Figs. 5-7) are summarized in

Fig. 13- Curves of each temperature extrapolate to a zero value of Ao/ASLne as 0 approaches zero. Thus, 0 ^

= 0 and the dislocation-dislocation interaction term dominates. This is similar to the single-crystal Ag behavior re- ported by Necking and Kocks^ and is probably typical of all relatively pure

fee metals. From Fig. 13 it also appears that m^ is temperature dependent,

increasing with increasing temperature. Deviation from linearity at higher

stress or temperature is associated with recovery, similar to that discussed

in the strain-hardening section.

114
Chr istodoulou and Jonas'*^ recently reported that their strain-rate

sensitivity parameter ( A2.no/AS,ne ) of oxygen-free 99.95 Cu, measured at room

temperature, initially decreased to a minimum and then began to rise. Other

inves t igat ions ,

^ ^ ^ some of which also obtained measurements at low

strains, do not support their observation. Additionally, our data of Fig. 13

clearly indicate a steady increase of Ao/A£ne over the entire deformation

range.

The strain-rate dependence of A1 at low temperatures is summarized in

Fig. 1^4. These results agree qualitatively with Takai et al.,^^ who measured

the strain-rate sensitivity of 1050 A1 (99.6 Al) at 4, 76, 120, and 295 K and

found a maximum of (Ao/A£ne)^ at 120 K. Very small (Ao/A2,n£)rp values at 4 K were also reported by Thornton et al.^

Whereas Ao/Mnt values of relatively pure Cu extrapolate to zero over

the entire 4-295 K temperature range, only the very low temperature (4-76 K)

Al values extrapolate to zero stress. Above 76 K the Al data appear to ex- trapolate to positive values. It is unlikely that the solute frictional

stress, Oj., is lower at lower temperatures . Typically one considers this

stress to originate from short-range stress fields and, thus, to decrease at

higher temperatures with assistance from thermal activation. If is positive, the term m^ - m^ must approach zero at low temperatures; thus, m^. must

approach zero as T approaches zero more quickly than m^ does (since the low

temperature data reflect dislocation Interaction behavior, not solute frictional stress behavior). At higher temperatures , the o^(m^ - m^) term gradually becomes dominant as the slope of Aa/Ailne decreases. The reduction of the

extrapolated values of Ao/AHne at zero stress at higher temperatures now reflect the temperature dependence of o^. This suggests a transition from a short-range, dislocation-interaction rate-controlling deformation mechanism

at low temperatures to a long-range, dislocation-solute interaction mechanism

at ambient temperatures. Decreases of extrapolated Aa/Afi,ne values (corresponding to 0 ^) with increases of test temperature at T ^ 100 K have also

been reported in solid-solution-strengthened Ni-base, Cr, Fe alloys, Ni-Co

alloys, and Al-Fe alloys.

Finally, we found a clear dependence of Ao on o (or e) that decreased at

both high and low temperatures. This dependence makes strain-rate sensitivity averaging techniques precarious. In measuring strain-rate sensitivity

as a function of temperature for Al-Fe, Bull and Saimoto^° averaged the

strains between a minimum of 0.04 and a maximum where 0=0. Two facts oppose

this practice: (1) The strain at which the strain hardening reaches zero is

dependent on a number of testing parameters; (2) Any averaging range based on strain range results in various stress ranges corresponding to each temperature, thus affecting the range of Aa data points to be included. It appears

best to compare Ao/AHne data versus temperature at constant stress (a) or at

constant normalized stress (a - Oy).

As discussed earlier, we observed transient yield point effects from

strain rate changes in both Cu and Al at temperatures above 76 K. In previous studies of transient effects from strain-rate changes , most re- searchers attributed the yield point effects to changes of mobile dislocation

115
density, but van den Brink et al.^3 made strong measurement correlations

between Ao^ changes and dynamic strain aging (see Fig. 3)- However, we observed larger values of Ao^ in Cu, the purer material. It is most likely that

our transient effects were caused by changes of dislocation density, not

dislocation mobility.

Activation Volume

Usually the activation volume parameter has been calculated from

strain-rate sensitivity measurements. This parameter is defined^"^ as

(9)

number and distribution of mobile dislocaactivation process controlling the rate of

e = exp (-Q/kT) , (10)

with Q the activation energy for the process. The parameter V is considered to be related to the average volume of the rate-controlling process by

V = bdJ, , (11)

with d the distance over which the dislocation interaction process proceeds

and I the dislocation interaction length.

Experimentally, it has been customary to consider to be independent

of 0 (Eq. 9) and thus to measure the term kT(A£ne/Ao)^ to obtain V. However,

one can easily obtain (see, e.g.. Frost and Ashby^°) a slightly modified rate

equation for obstacle-controlled deformation by using the following two relationships: The strain rate is related to dislocation density and average

velocity (v) by

e = Pm^v , (12)

and is related to o by

0 = a pb (13)

where a, a constant, is near unity.

The average dislocation velocity may be represented by the rate

equation

V = kT

3 Znc/c^

3 0 where is a function of the

tions. This assumes a single

deformation.

V = Bbv exp [-AG(o)/(kT)] , (14)

where B is a constant, v is the frequency, and AG is the Gibbs free energy

for overcoming obstacles.

116
Combining Eqs. 12 through yields

e= Bv(o/au)^ exp [~AG( o ) / (kT ) ] , (15)

and we see that the pre-exponential term does include stress. Differentiating Eq. 15 with respect to a,

kT

_

Kocks et al.^^ and Frost and Ashby^® describe AG as

AG(o) = AF [1 - a/o^] (17)

for relatively large values of AF, where AF is the total free energy needed

to overcome obstacles in the absence of applied stress. The constant is

described as the athermal flow strength or, normally, the yield strength at

absolute zero. Here,

3 Une

j do

3AG(o)

3 a

( 16 )

V = AF/Oq , (18)

when AF is assumed to be large, of the order of 1 ub^. On combining

Eqs. 16-18,

kT

3 ( AF • o/ Oq ) 3o

V

(19)

The contribution of the pre-exponential term in Eq. 19 for our Cu and A1

data is about of A£ne/Ao. Therefore, in our data presentations

this term has not been included; this permits better comparison with past

measurements. We subsequently refer to this quantity, kT ( A2,ne/Ao) , as the

apparent activation volume, V*.

In Fig. 15 the apparent activation volume, V*, of Cu is plotted versus

test temperature at constant stress. All data, although approaching close to

V* = 0 as o^O, clearly extrapolate to small positive values of V* as o-»-0. Butt and Feltham^^ showed that £nV* is linearly related to Una at constant T

for oxygen-free copper; this is substantiated by our data (plotted in this manner in Fig. 16). The slope of the curves at 23 K and above are approximately equivalent; however at 4 K this equivalence does not hold. Unlike Butt

and Feltham, we did not find that the slope of UnV* versus 8,no was -1 (oV* = constant at a given temperature), but instead was less than -1 at temperatures ^ 23 K. However, at 4 K the slope of the curve plotted in Fig. 16 is

about -1

.

If one assumes a physical significance in the absolute values of V*, we realize that these values are quite large. Interactions of dislocations

within volumes from 100 to 1000 b^ indicate quite long distances between interaction points. If d = b, then i is in the range 100 to 1000b. These

117
large activation volumes for Cu are consistent with all

have calculated v* . ^ ^ ^ ^ ^ ^ ^ ^

previous reports that

For Al, the values of V* are plotted versus temperature at constant a in

Fig. 17- When one considers that Oy of Al is almost equivalent to that of

Cu, the values of M* for Al are less than those of Cu. Secondly, at very low

temperature (4 K), the V* values are larger than those at 20 K. Similar to

that of Cu, the 5,nV* versus 2,no of Al is linear at low temperatures, as shown

in Fig. 18. At higher temperatures (2 150 K), slight curvature of the logarithmic dependence is apparent.

Saimoto et al.^^~^^ reported activation volumes ([1/T] [ 9iino/3 S,ne]

,

which they label "strain-rate sensitivities") for Al-dilute Fe alloys over

the temperature range 41 to 41 8 K. They found a minimum of V but at a much

higher temperature, near 200 K. As mentioned earlier, their A£na/AJlne was averaged over a range of stresses, whereas we plotted Ao/AHne at a constant

stress level. Like Basinski,^^ they found no clear trend below about 77 K

(both used the "averaging" method of strain-rate sensitivity analysis).

There have been many reports of dilute alloy softening at very low temperatures. Decreases of yield strength of Cu- and Ag-based dilute alloys

have been found at temperatures less than about 50 K. The minimum of V* in

dilute Al-Fe at about 20 K, opposed to the lack of a minimum in 99.99 Cu,

may be a sensitive indicator of this phenomena.

Another possible explanation for the apparent softening associated

with the increase of V* at very low temperatures is adiabatic heating. This

should manifest itself with a strong dependence on e and will be explored in

the near future.

CONCLUSIONS

A study of the strain-hardening and strain-rate sensitivity of 99.99 Cu

and 99.1 Al over the temperature range 4 to 295 K found the following ratecontrolling deformation features:

The strain-hardening characteristics are best portrayed using the term

(a - Oy) (da/de) to describe dislocation accumulation as a function of normalized

applied stress, o - Oy. Two regions are found in Cu and Al : ( 1 ) an athermal,

linear range of strong increase of mobile dislocation density and (2) a range

associated with dynamic recovery, where the rate of increase of dislocations

levels off and then decreases.

The strain-rate sensitivity characteristics were identified in two ways

(1) as a strain-rate sensitivity parameter, ( Ao/Afi-ne ),p, and (2) as an activation volume, kT ( A2,ne/ Ao )^ . The strain-rate sensitivity of Cu indicates that

the frictional stress between dislocations and solute is effectively zero for

99.99 Cu, but positive for 99.1 Al. This is also indicated in the increased

temperature dependence of the yield strength of the Al. The effect of solutedislocation interactions may be reflected in the low temperature minimum

of the activation volume, since it is absent in the Cu.

118
ACKNOWLEDGMENT

Partial support for this project was kindly supplied by the Office of

Fusion Energy, DOE.

NOTATION

b Burgers vector

B constant

c solute content

d distance affected by dislocation-interaction process

AF total free energy (needed to overcome obstacles in the absence

of applied stress)

AG Gibbs free energy

k Boltzmann's constant

Z length of dislocation interaction

rrid /dlno^\

\31ne

m

^

/ 31no^\

\31ne /-p

p active dislocation density

Q activation energy

T temperature

V average dislocation velocity

V activation volume

V* apparent activation volume

^ * ^

a constant

£ tensile strain

•£ strain rate

119
Cq strain rate depending on number and distribution of mobile

dislocations

0 strain-hardening rate

0|^ athermal strain-hardening rate

0p strain-hardening rate related to dynamic recovery

u shear modulus

V frequency

mobile dislocation density

a tensile flow stress; when used in strain-rate change measurements, it represents the average between initial (low e) and

maximum (high e) flow stress.

Oq constant = athermal flow strength or yield strength at 0 K

0 ^ dislocation-dislocation interaction contribution to flow strength

0 ^ frictional strength contribution to flow strength due to dislocation-solute interactions

Aa change in flow stress resulting from increase in strain rate

measured by comparison of initial stress (low e) to maximum

stress (high e)

.

Ao^ change in flow stress resulting from increase in strain rate

measured by comparison of steady-state conditions (see Fig. 3a)

da/de strain hardening

(do/lne)rp strain-rate sensitivity

REFERENCES

1. H. Conrad, in "High Strength Materials," Wiley & Sons, New York (1965),

p. J436 .

2. K. A. Warren and R. P. Reed, "Tensile and Impact Properties of Selected

Materials from 20 to 300 K." NBS Monograph 63, U.S. Government Printing Office, Washington, D.C. (1963).

3. U. F. Kocks, R. E. Cook, and R. A. Mulford, Acta Metall. 623 (1985).

4. H. Mecking and U. F. Kocks, Acta Metall. 1865 ( 1981).

5. R. A. Mulford and U. F. Kocks, Acta Metall. 27, 1125 (1979).

120
. T. A. Bloom, U. F. Kocks, and P. Nash, Acta Metall. 265 (1985).

7. R. A. Mulford and U. F. Kocks, Scripts Metall. J_3, 729 ( 1979).

8. R. A. Mulford, Acta Metall. 27, 1115 (1979).

9. P. R. Thornton, T. E. Mitchell, and P. B. Hirsch, Philos. Mag. 7, 337

( 1963) .

10. R

.

Zeyfang, 0. Buck, and A. Seeger, Phys. Status Solid! (b)6l , 551 ( 1974)

11. P. C. J. Gallagher, Philos. Mag. 11, 51 (1967).

12. H . Neuhauser, N. Himstodt, and Ch. Schwink, Phys. Status Solid! 93, 585

(1970).

13 . H . Conrad, Acta Metall. 6, 339 (1958).

14. N . Chr istodoulou and J. J. Jonas, Acta Metall. 32, 1655 (1984).

15 . M. Z. Butt and P. Feltham, Metall. Sci. 18, 123 (1984).

16. G. J. den Otter and A. van den Beukel, Phys. Status Solid! (a)55, 785

(1979).

17. G. Regazzoni and F. Montheillet, Institute of Physics Conferer.ci Series

No. 70, The Institute of Physics, London (1984), p. 63.

00 > Melander and A. Thuvander, Res. Mech. 5, 129 (1982).

ISl S. Basinski, Philos. Mag. 3, 393 ( 1959').

20. T. Takai, T. Tatsumi , and T. Tamura, "Strength of Metals and Alloys,"

Pergamon Press, New York (1985).

21 . S. K. Mitra, P. W. Osborne, and J. E. Dorn, Trans. AIME 221, 1206

(1961 ).

22. S. Saimoto, H. Sang, and L. R. Morris, Acta Metall. 29, 215 (1981).

23. s. Saimoto and M. S. Duesberry, Acta Metall. 32, 147 (1984).

24. S. Saimoto and H. Sang, Acta Metall. 31, 1873 (1983).

25. M. J. Bull and S. Saimoto, "Strength of Metals and Alloys," Pergamon

Press, New York (1985), p. 893.

26. M. J. Bull and S. Saimoto, "Proc. International Solute - Defect Interaction: Theory and Experiment," March (1986).

27. S. K. Mitra and J. E. Dorn, Trans. AIME 227, 1015 (1963).

121
28. C. Perdrix and F. Montheillet, "Strength of Metals and Alloys," Pergamon

Press, New York (1985), p. 571.

29. B. A. Parker, "Strength of Metals and Alloys," Pergamon Press, New York

(1985), p. 899.

30. F. P. Bullen and M. M. Hutchinson, Philos. Mag. 7, 557 (1962).

31. D. T. Read and R. L. Tobler, "Advances in Cryogenic Engineering

Materials" vol. 28, Plenum Press, New York (1982), p. 17.

32. R. P. Reed and J. M. Arvidson, "Advances in Cryogenic Engineering

Materials," vol. 30, Plenum Press, New York (198A), p. 105.

33. S. H. van den Brink, A. van den Beukel, and P. G. McCormick, Phys.

Status Solidi (a)30 , ^69 (1975).

3^. B. Wielke and G. Schoeck, Phys. Status Solidi (a ) 38

,

539 (1976).

35. B. Wielke, Acta Metall. 103 ( 1978).

36. F. Frohlich and P. Gran, Phys. Status Solidi (a )

1 A

,

1 15 ( 1972).

37. G. Schoeck, Phys. Status Solidi 8, H99 (1965).

38. H. J. Frost and M. F. Ashby, Deformation Mechanism Maps: The Plasticity

and Creep of Metals and Ceramics, Pergamon Press, Oxford, England

(1982), p. 6.

39. U. F. Kocks, A. S. Arogn, and M. F. Ashby, Prog. Mater. Sci. 19 (1975).

Table 1 . Material Characterization

Alloy Chemical Hardness Grain

Composition (Rg) Size

( wt . 5?

)

(pm)

CDA 101 Cu 99.99 Cu + Ag 22 36

1100 A1 99.1 A1 + 0.6 Fe, 0.2 Cu, 0.1 SI 36 1 45

122
LN2^LHe Stainless Steel

Fig. 1. Cryostat used for conducting strain-rate sensitivity

tests at temperatures from 4 to 295 K.

123
a. Effects of slackening of crosshead

following change of strain rate from

high to low

b. Effects of low time constant for

filtering of load signal

Fig. 2. Measurement variables affecting stress-strain

dependence following strain-rate changes.

124
a Transient at high strain rate

.

b. Transient at Low strain rate

c. Rate of strain hardening

dependent on strain rate

at low strains in the

high strain-rate region

d. Discontinuous yielding

(adiabatic heating) at 4 K

e. Incursion of discontinuous

yielding on strain rate

change at T ^ ~20 K

Fig. 3. Variables affecting stress-strain characteristics

following strain-rate changes.

125
Fig, 4. The tensile stress-strain curves of CDA 101 Cu at various temperatures.

126
8

CDA 101 Copper

o 295 K

O)CO

UJ

cc

4 —

<S) <3 3

2 —

OolS

«»igP“ a O

iP“ o CO

oO

corf’

Oo'Srf'*' "

1°

a

20 40 60 80 100 120 140 160 180 200 220

FLOW STRESS, MPa

Z40

Fig. 5. Stress change of CDA lOL Cu due to increasing the strain rate

(2 X 10“^ to 2 X 10“^ s”^) as a function of applied stress at

room temperature. Data are from three independent tests.

260

127
10

9

7

CDA 101 Copper

o 23 K A I50K

X 200 K

• 250 K

o

CL

CO

LlI 5

CL

Ico

0

• A

• X ^

X A

V AA

.XA ^

AA

•;oa5o!

• • • • aX •• AyX A^AAA^^ O

O O

^ O X O O O

° o

A 6o Ox d O O

Ov X

o o

50 100 150 200

FLOW STRESS, MPa

250

Fig. 6. Stress change of CDA 101 Cu due to increasing the strain rate

(2 X 10“^ to 2 X 10“3 s"M as a function of applied stress at

various test temperatures.

300

128
10

9

8

7

CDA 101 Copper

o 4 K

• 50 K A 76 K

• 98 K

o

9r 6

to

to 5

Ld

octc/) 4

<1

A A A

* .A OOOc|oO OpPOOtj^ OqO 0°|3 o o o loo

50 100 150 200

FLOW STRESS, MPa

250 300

Fig. 7. Stress change of CDA 101 Cu due to increasing the strain rate

(2 X 10“^ to 2 X 10”^ s”^) as a function of applied stress at

various test temperatures.

350

129
STRESS,

MPa

130
A

STRESS,

MPo

Fig. 9. Stress change of 1100 Al due to

(2. X 10-5 to 2 X 10"5 s"^) as a at various test temperatures.

increasing the strain rate

function of applied stress

131
Fig. 10. Stress change of 1100 Al due to increasing the strain rate

(2 X 10“^ to 2 X 10“^ s“^) as a function of applied stress

at various test temperatures.

132
.

10'

MPa

200

b

<1 '

b

I ^b

150

100

50

+Q

+ O CDA 101 Copper

4K

• 50 K

+ 76K

o 198 K o 295 K

50 100 150

cr-CTy , MPa

200 250 300

Fig. 11. Strain-hardening characteristics of CDA 101 Cu at various temperatures

133
10

MP

Fig. 12. Strain-hardening characteristics of 1100 A1 at various temperatures.

134
A

Inc

FLOW STRESS, MPa

Fig. 13. General characterization for CDA 101 Cu of the effects of s rate change (2 x 10“ to 2 x 10“ s“^) on the stress change

function of applied stress at various test temperatures.

trainas a

135
MPa

FLOW STRESS, MPa

Fig.^14. General characterization for 1100 Al of the effects of strain-rate c » - T change (2 x 10“^ to 2 x 10“'^ s" ) on the stress change as a function

of applied stress at various temperatures.

136
ACTIVATION

VOLUME,

V

Fig. 15. Apparent activation volume versus temperature

at constant stress levels for CDA 101 Cu

.

137
Fig. 16. Apparent activation volume versus stress

at constant temperature for CDA 101 Cu

.

138
800

1 100 Aluminum

Fig. 17. Apparent activation volume versus temperature

at constant stress levels for 1100 Al

.

300

139
FLOW STRESS, MPa

Fig. 18. Apparent activation volume versus stress

at constant temperature for 1100 Al

.

lAO
w, -.,

t^sy r K -'^ '^,5- ’ *

y t ' f.- i

»

^ Xg 4 i>*M i f ^fcayi

Mirit4i0nn«ri'4^^ :ii^ft<iliv»l»;«»'^ '

[rhlij^r tVtaff Wt^ *'• ,*. '? '* »v V' y

’ ^W'-.-

' "' "• '..!'•'•

• £ “

l' . b-T . , '-fc KE «, ’- • '

'j' t iT «s"

‘'i.

*

' — " K'.'

' **

*'' ' * '

jk

EFFECT OF PRIOR DEFORMATION ON THE 76“K FRACTURE TOUGHNESS

OF AISI 304L AND AWS 3 O 8 L STAINLESS STEELS*

P. T. Purtscher and D. T. Read

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

AISI 304L and its weld metal, AWS 308 L, may undergo a partial transformation to martensite during cryogenic service owing to thermal and mechanical

stresses. In this study, the effect of service-induced deformation on the

toughness of these materials was determined. Low temperature compressive

loading in the laboratory produced larger deformations. Crack initiation

toughness, Kj^(J), and tearing resistance, dJ/da, at 76 K were evaluated as a function of martensite content, a measure of the deformation in these steels.

The results showed that the toughness properties of the 304L decrease gradually as the martensite content increases from the 5 to 85^ level found in the

service condition to the 455^ level obtained by compressive loading. The

decrease was less than that expected on the basis of the increased flow

stress. The toughness properties of the 3 O 8 L weld metal decreased more

sharply with increased martensite content than those of the 304L. The sharp

decrease is associated with a degradation of the properties of the delta

ferrite rather than that of the austenite.

Key words: austenitic

mater ials.

stainless steel, cold work, cryogenic properties of

delta ferrite, fracture toughness, martensite

*Contr ibution of NBS; not subject to copyright.

141
INTRODUCTION

Austenitic stainless steels are commonly used for cryogenic structural

components because they have excellent low temperature properties, specifically, high elastic modulus, high toughness and ductility, and good weldability. Alloy AISI 304L, which contains nominally 18 wt.^ Or, 8 wt.^ Ni,

and less than 0.04 wt.^ C, is commonly used with the filler material AWS

308L, which contains approximately 2 % more chromium and nickel.

Under cryogenic operating conditions, these structural alloys may partially transform to martensite owing to a combination of thermal effects

and plastic deformation resulting from thermal or mechanical stresses. Two

types of martensite form in these steels: e, a nonmagnetic phase, and a', a magnetic phase. Magnetic measurements provide an easy way to monitor the

a' martensite content in these steels. Ogawa and Morris^ showed that prior

deformation improved the relationship between strength and toughness for a Fe-Mn-Cr austenitic alloy at 4 K, but only e martensite forms in this alloy.

The a’ martensite may degrade the toughness of the alloy significantly

because it has the same bcc crystal structure as delta ferrite, which promotes brittle fracture in austenitic stainless steel castings. “ Whenever

martensite is mentioned in the rest of the text, it refers to the a' phase.

Austenite-to-martensite transformation occurs when steels are cooled

below their martensite start temperature or when they are plastically deformed below their martensite deformation start temperature. When straininduced martensite forms in areas of plastically deformed austenite, the

toughness of the material remains relatively high. The effect of martensite

in the microstructure prior to testing has not been studied systematically.

A thorough discussion of martensitic transformation in austenitic steels was presented by Reed.^

In this study, we evaluated the toughness of 304L and 308L steel obtained from a component that had been in cryogenic service. We also studied

the effect of larger deformations. We used a simple fracture mechanics model

to illustrate the effect of increased deformation and the resulting increase

of the martensite content on the critical crack size required for the initiation of crack growth.

MATERIALS

Specimens of AISI 304L were removed from a large, forged flange that was

part of a liquid natural gas facility, where it experienced temperatures as

low as 1 1 1 K. Test pieces, 115 mm long and 25 mm square, were machined so that the long direction would be parallel to the forging flow lines (similar

to an L-T orientation).

Weld metal specimens of AISI 308L were removed from a pipe that joined

the large, forged flange. The specimens, 16 mm square and 115 mm long, were machined from the pipe-to-f lange connection so that the weld metal was in the

center of the bar.

142
Round tensile specimens of 304L and 308L were machined to a 50-mm gauge

length and 6-mm diameter. In the flange, the tensile bars were oriented

along the forging flow lines. In the weld, the tensile bar was machined

transverse to the length of the pipe so that the tensile bar was completely

weld metal.

A commercially supplied, spiral spring-balanced, magnetic measuring

instrument was used to determine the magnetic susceptibility of the materials.

The instrument was calibrated according to AWS A4.2~7^ standard procedure.

The 304L specimens in the as-received, service condition had a magnetization

that corresponded to a ferrite number (FN) of 6 to 9 . The 308 L weld metal

specimens had a magnetization of approximately 8 FN.

MAGNETIC SUSCEPTIBILITY RESULTS

Initial attempts to induce additional deformation by thermally cycling

the as-received 304L and 308 L material proved fruitless. A small piece of

each material was cooled from room temperature to 76 K and then back again

five times without any increase in magnetization. This short experiment

proved that the materials in the service condition were thermally stable with

respect to the martensitic transformation.

In the next step, the pieces were heat-treated at 1060°C for half an hour in an argon atmosphere and water quenched. Then the annealed pieces

were thermally cycled from room temperature to 76 K eight times. This procedure produced 3 to martensite in the 304L and 0.5??^ martensite in the 3 O 8L

before the magnetization appeared to saturate. The values were less than the

magnetization of the as-received materials. Additional thermal cycling from room temperature to 4 K did not produce any further change in the martensite

content. This experiment indicated that the two annealed materials were not

completely stable with respect to the martensitic transformation and that

the martensite found in the service condition was not due solely to thermal

effects.

THERMOMECHANICAL TREATMENTS

Selected specimens of the 304L and 308 L material were heat-treated in an argon atmosphere at 1060°C for 30 min. and then water quenched. This annealing treatment did not change the grain size appreciably, which was approximately 100 ym in the 304L, but it did remove all the martensite that formed

during the time in service at 1 1 1 K. After heat treatment, the 304L had a FN

of 1, and the 308 L weld metal had a FN of 4. The residual magnetic susceptibility in the steels after heat treatment was due to the presence of delta

ferrite.

Several specimens of the 304L and one 3 O 8 L specimen were subjected to

compression loading to increase the cold work in the microstructure . The

specimen to be compressed and two larger pieces of steel were cooled to

liquid nitrogen temperature by immersion. Then the three pieces were removed

from the liquid nitrogen; the specimen was placed between the two larger

143
pieces of steel to insulate it from the surroundings. A servohydraul ic test

machine with a 1 -MN load capacity applied the compressive load through the

buffer plates. Different loads were used to produce different levels of

strain-induced martensite.

TOUGHNESS TESTING PROCEDURE

Selected specimens were heat-treated, loaded in compression, and then

notched to produce bend specimens. The fatigue precracking and subsequent

testing were performed in a three-point bend fixture with a span that was

four times the width of the specimen. The specimens were maintained in a liquid nitrogen bath, 76 K, during all the mechanical testing. A servo- hydraulic test machine with a load capacity of 100 kN was used.

During the toughness tests, two clip gauges for cryogenic service were

employed, one to monitor the mouth-opening displacement of the bend specimen

and a second to measure the load-line displacement of the specimen with

respect to a comparison bar. From the output of the mouth-opening displacement and the load, the initial crack length and subsequent crack growth were calculated during the test by the unloading compliance technique. ^ From

the load and load-line displacement information, values of the J-integral

were calculated according to ASTM E 8l3~8l test procedure."^ A computeraided test analysis technique was used, similar to that used for compact

tension specimens;® both techniques were developed at NBS-Boulder to simplify

and standardize the test procedure.

The results of each test were summarized by a plot of the J-integral vs. the change in crack length during the test (J-R curve).’ The crack lengths

were corrected after the test by measuring the fatigue precrack and the final

crack length and using linear interpolation to scale each intermediate crack

length. The J-integral values were recalculated for each unloading point

using Ernst and Paris's formula^ and the corrected crack lengths.

The Jj^ fracture toughness value for each specimen was determined by

the intersection of the calculated blunting line* with the measured data

points, as specified by ASTM E 8l 3~8l

.

The value of Jj^ refers to the energy

required to initiate tearing. In addition, a value of dJ/da, the tearing slope

over the first 1.5 mm of tearing, was calculated for each specimen. This

value refers to the resistance of the steel to crack growth as a function of

martensite content.

The correction technique described above produces reliable J-R curves as long as the actual final crack length agrees reasonably well with the final

crack length measured by the compliance technique and the correction factor

is close to 1.0. If the correction factor is greater than 2, past experience

indicates that the dJ/da values are not a reliable description of the crack

growth. The value of Jj^ measured in the test is unaffected by this correction factor.

*The blunting line is defined as J = 2a^j_Q^Aa, where o^^ow average of

the yield and ultimate strength; a value of 690 MPa was used for these tests.

144
Estimates of the plane strain fracture toughness parameter, Kj^, were obtained from the relationship:

Ki^(J) = (1)

where E is Young's modulus. The value of E for stainless steels at 76 K was estimated to be 206.8 GPa.^^

MECHANICAL TESTING RESULTS

The results of the tensile testing show that both 304L and 308L have

relatively low initial flow strengths, low work -harden ing rates, and good

ductility. Table 1 contains a summary of the tensile properties. Magnetic

susceptibility readings were taken from the deformed gauge length after

testing. During the test, the 304L transformed to approximately 83?

martensite, and the weld metal, to about 69? martensite.

Figure 1 shows representative load vs. mouth -open ing displacement curves

from the toughness tests for both 304L and 308L. In all the tests, was

found at a point on the curve near the maximum load. An approximation of

can be obtained by estimating the area under the curve up to the maximum

load. An indication of the relative value of dJ/da can be obtained by

observing the behavior of the curves after the maximum load is reached. If

the load drops off very slowly, dJ/da will be high, corresponding to a high

resistance to tearing. If the load decreases rapidly, dJ/da will be low,

corresponding to a lower resistance to tearing. The load-displacement curves

in Figure 1 indicate that the toughness was reduced by additional deformation, shown by the presence of more martensite in the microstructure. The

effect is particularly noticeable for the 308L weld metal.

Values of and dJ/da were calculated for each of the specimens and

are listed in Table 2 along with the average level of martensite present

before testing. The values of Kj (J) are displayed graphically in Figures 2

and 3.

No 304L annealed specimen was tested because of the scarcity of material. The difference between 304 and 304L is anticipated to be relatively

small, with 304L having a slightly higher Kj^(J). A previously reported data

point (•) for AISI 304 25-mm-thlck plate material in the annealed condition”'^

is shown in Figure 2 for comparison. In addition, two unpublished values (A)

of Kj^ obtained from a 6-mm-thick sheet of AISI 304 with larger deformations,

higher martensite contents,”'^ are also shown to define the trend line more

clearly.

The toughness of the 308L specimen in the service condition was significantly lower (34?) than the toughness of the annealed specimen, as shown in

Figure 3. A similar decrease can be seen for the toughness of the 304L

specimen in the service condition compared with that reported for 304 in the

annealed condition.

The effects of larger deformations can be evaluated from the results for

the specimens in the service-and-de formed condition. In the base metal.

145
304L, Kjq('J) di"ops off rather slowly and the tearing slope does not change

significantly (Figure 2). In the 308L weld metal, Kj^(J) decreases significantly as the martensite content increases (Figure 3).

The J-test fracture surfaces of the specimens were studied with a scan- ning electron microscope (SEM) . No significant differences between the

service condition and service-and-deformed specimens of 304L were apparent;

all exhibited dimpled rupture features. Figure 4 shows representative photomicrographs of the surface of the specimen with martensite.

The fracture surfaces of all three weld metal specimens were different.

The fracture surface of the annealed specimen was composed of extremely fine

dimples with only an occasional flat area (see Figure 5). Energy dispersive

X-ray analysis (EDAX) showed that the flat areas had a higher chromium content and lower nickel content than the dimpled area, indicating that the flat areas are regions of delta ferrite in the austenitic matrix.

The surface of the service condition specimen is shown in Figure 6. A higher percentage of the fracture area is composed of the flat areas as compared to the fracture surface of the annealed specimen. The fracture surface

of the service-and-deformed specimen (Figure 7), is almost completely devoid

of the dimpled rupture features.

DISCUSSION

The test results indicate that prior cold work, as shown by the increase

in the martensite content from 0 to 455^, can affect crack initiation toughness, Kj-^(J), and tearing resistance, dJ/da, of 30AL and 308L specimens. The

effect of prior cold work was more pronounced in the 308L weld metal that

contained 4^ delta ferrite.

In evaluating these test results, the effect of yield strength must be

taken into account. The deformation at 76 K increases the yield strength by

two different mechanisms: martensitic transformation and strain hardening.

An austenitic stainless steel, such as AISl 310, that is more stable with

respect to the martensite transformation than 304L, would suffer a loss of

toughness after deformation at 76 K without the formation of any martensite.

The loss of toughness in this case is associated with the increase in yield

strength by strain hardening.

In this set of tests, the combined effect of the martensitic transformation and strain hardening was measured. Because some specimens were deformed

in compression first and then tested in tension, an accurate estimate of the

tensile yield strength was not possible and the two effects could not be

separated. For the purposes of this study, analyzing the effect of prior

deformation on the toughness of 304L and 308L, it is appropriate to consider

the combined effects for engineering applications because plastic deformation

always transforms some austenite to martensite. However, these results and

conclusions are valid only for metastable alloys like 304L and 308L and

should not be applied to other alloys without review.

146
It is possible to estimate the flow stress for each specimen from the

loads used in compressing the specimens in the laboratory. From these estimated values, a decrease in Kj^(J) that would be predicted for annealed

austenitic steels at 4 K can be calculated.^ Table 3 summarizes the results

of our estimated flow stresses and the effect that the change in flow

strength has on the toughness. The flow stresses in tension would be less

than the estimated values shown, but the estimates represent a reasonable

measure of the relative increase in flow stress with increased cold work and

the effect of increased strength on Kj^(J). The predicted loss in toughness

for the base metal is more than the measured change in toughness. With

increases in strength, the toughness of the base metal apparently decreases

at a lower rate than predicted for these cold-worked, metastable steels. For

the weld metal, the opposite is true; the predicted -decrease in toughness is

less than the measured decrease in toughness.

The results raise a question concerning the difference between base and

weld metal. The SEM examination of the base metal specimens shows that mar- tensite does not di-rectly affect the fracture mechanisms (Figure 4); that is,

the fracture morphology of the service condition specimen with 5 to 8% mar- tensite is not appreciably different from the fracture morphology of the

specimen with 45(^ martensite: there is no ductile-to-brittle transition. A

304L specimen with 0^ martensite was not tested; however, previous experience

indicates that the fracture surface features of a 304L specimen without mar- tensite would not be significantly different from those of the other 304L

specimens.

For the weld metal specimens, prior deformation does produce changes in

the fracture morphology (Figures 5, 6, and 7). The properties of the ferrite

in the weld metal microstructure are apparently degraded by the cold deformation more than the properties of the austenite, and as a result, the crack

prefers to grow through the delta ferrite.

Fracture Mechanics Considerations

Fracture mechanics provides a means for assessing the fitness of a particular material for a given application. A simple analysis involves a small through-thickness crack in the center of a plate with finite thickness

and other dimensions large with respect to the length of the crack, a. The

driving force for fracture is given by

K = o(TTa/2)^^2 (2)

where K is the applied stress intensity factor and a is the applied stress.

Different formulas are required for different geometries, stress concentrations, or residual stresses. The initiation of tearing is predicted when K exceeds Kj^(J).

For the 304L in the service condition,

(Table 2). If the stress on the plate were

before the predicted initiation of tearing,

base metal specimen, Kj^(J) was 157 MPa/m.

Kic(J) was about 200 MPa/m

100 MPa (50% of yield), then a For the most severely deformed

If the stress were 100 MPa, the

147
critical crack size for a center-cracked plate would be 1.5 m. In the 308L

weld metal, Ky^(J) was 18 M MPa/m in the service condition. If the stress were 122 MPa (half of yield), the critical crack size would be about 1.5 m.

In the worst case,_for the serv ice-and-de formed weld metal specimen

where was 95 MPa/m, if the stress were 122 MPa, then the critical

crack length would still be 0.4 m. The size of the critical crack in each of

the aforementioned situations indicates that the material possesses good

toughness with respect to crack initiation. Specific structural situations,

especially those Involving high stresses, stress concentrations, and residual

stresses, require more detailed analysis. A small factor of safety is built

into the toughness values because the test temperature was lower than the

actual service temperature.

In summary, prior deformation decreases the crack tolerance of these

materials, but large deformations are required for significant effects.

These initial tests indicate that problems are more likely to be found in the

weld than the base material.

CONCLUSIONS

For both 304L and 308L, cryogenic service resulted in lower

values than those reported for the annealed condition. Low temperature com- pressive loading further degraded the resistance to crack initiation and

tearing; the 308 L weld metal was affected more severely than the 304L base

metal.

These results do not permit the effects of individual microstructural

components, such as strain-induced and thermally induced martensite and delta

ferrite, to be separated from the strain-hardening effects. However, it

appears that the presence of strain-induced martensite in the microstructure

is not detrimental to the fracture toughness properties.

ACKNOWLEDGMENTS

The authors thank the Mobil Oil Corporation for supplying the materials

and financial support to this work and the Department of Energy who supported

the development of the clip gauges and test procedures used in this study.

We also thank Frank Furillo of Mobil and David McColskey of the Fracture and

Deformation Division of NBS-Boulder for their help.

REFERENCES

1. R. Ogawa and J. W. Morris, Jr., "Influence of Processing on the

Cryogenic Mechanical Properties of High Strength High Manganese Stainless

Steel," in Advances in Cryogenic Engineering--Material s, Vol. 30, A. F.

Clark and R. P. Reed, eds. , Plenum Press, New York, 1984, pp. 177-184.

2. E. L. Brown, T. A. Whipple, and R. L. Tobler, "Fracture Toughness of CF8

Castings at Four Kelvin," Metallurgical Transactions , Vol. 1 4A , No. 8,

1983, pp. 1179-1 188.

148
3. R. P. Reed, "Martensitic Phase Transformations" in Materials at Low

Temperatures , R. P. Reed and A. F. Clark, eds. , American Society for

Metals, Metals Park, Ohio, 1983, PP- 295~3^1

•

, "Calibrating Magnetic Instruments to Measure the Delta Ferrite Content

of Austenitic Stainless Steel Weld Metal," AWS A4.2-74, American Welding

Society, Miami, Florida, 197^.

5. R. J. Buzzard and D. M. Fisher, "Load Displacement Measurements and Work

Determination in Three-Point Bend Tests of Notched and Precracked Specimens," Journal of Testing and Evaluation , Vol. 6, No. 1, January 1978,

PP. 35-39.

6. J . G. Merkle, R. G. Berggren, and W. J. Stelzman, "Compliance and Fracture," HSST Quarterly Progress Report , ORNL Report TM-il805, Oak Ridge

National Laboratory, Oak Ridge, Tennessee, March 1975.

7. ^ Measure of Fracture Toughness" ASTM E 8l 3~8l , in Annual Book of

ASTM Standards , Part 10, American Society for Testing and Materials,

Philadelphia, 1981, pp. 763~78l

.

8. R. L. Tobler, D. T. Read, and R. P. Reed, "Strength/Toughness Relationship for Interstitially Strengthened AISI 30i1 Stainless Steels at 4 K

Temperature, " in Fracture Mechanics: Thirteenth Conference , ASTM STP

7^3, Richard Roberts, Ed., American Society for Testing and Materials,

Philadelphia, 1981, pp. 250-268.

9. H. A. Ernst and P. C. Paris, "Techniques of Analysis of Load-Displacement

Records by J-Integral Methods," Report NUREG/CR-1 22 , U.S. Nuclear

Regulatory Commission, January 1980.

10. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, "Elastic-Constant

Variability in Stainless Steel 304," Journal of Applied Physics , Vol. 51 , 1980, pp. 305- 309 .

11. D. T. Read and R. P. Reed, "Fracture and Strength Properties of Selected

Austenitic Stainless Steels at Cryogenic Temperature," in Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures--!! , F. R. Fickett and R. P. Reed, eds., NBSIR 79-1609, National Bureau of

Standards, Boulder, Colorado, 1979, pp. 81-122.

12. P. T. Purtscher, Fracture and Deformation Division, National Bureau of

Standards, Boulder, Colorado, unpublished data.

149
Table 1 . Results of Tensile Tests at 76 K

Yield Strength

(MPa)

Ultimate Strength

(MPa)

Reduct ion

of Area {%)

AISI 304L

annealed

16A (23.8 ksi) 1475 (214 ksi) 65

AISI 30ML

service condition

203 (29.5 ksi) 1450 (211 ksi) 68.5

AISI 308L

service condition

244 (35.4 ksi) 1360 (197 ksi) 69.5

Table 2. Calculated Toughness and Tearing Resistance

Condition Kic(J) dJ/da Martensite

( MPa/m) ( kN/m/mm) i%)

AISI 304L

Serv ice 200 * 5-8

Serv ice 208 166 5-8

Annealed and deformed 191 1 34 13

Service and deformed 190 195 20

Service and deformed 191 1 74 18

Service and deformed 183 * 35

Service and deformed 157 1 48 45

AISI 308L

Annealed 279 ~ 0 0

Serv ice 184 87 4 Service and deformed 95 25 19

* Not available

150
Table 3- Estimated Flow Strengths and Related Toughness Changes

Cond it ion Estimated Flow

Strengths (MPa)

Predicted Toughness

Change (MPa/m)*

Measured Toughness

Change (MPa/m)

AISI 304L

Serv ice 690 0 0

Serv ice 690 0 0

Annealed and

de formed

690 0 -1 0

Service and

de formed

860 -51 -10

Service and

deformed

860 -51 -10

Service and

deformed

930 -72 -1 7

Service and

de formed

AWS 308L

1 1 00 -1 23 -43

Annealed 690 0 0

Serv ice 690 0 -93

Service and

de formed

1035 -104 -184

^Calculated from reference 8, where Kj^(J) = 500 - 0.3(a) for annealed

austenitic stainless steels at 4 K. The o term is equal to the yield

strength.

151
LOAD

LOAD

5-8% a' 45% a'

p

max

DISPLACEMENT DISPLACEMENT

Figure 1. Load vs. mouth-opening displacement curves for (a) 304L in

service condition, (b) 30^L in service-and-deformed condition

with 45/S martensite, (c) 308L in annealed condition, and (d) 308L

in service-and-deformed condition with 195S martensite.

152
FRACTURE

TOUGHNESS.

K,^

(J).

MPaVnrT 400

? 300

200

100

PERCENT MAGNETIZATION

20 40 60 80

T

304

A Annealed Condition

• Annealded and Cold

Deformed

T“ ——r- 1 —

r

304L

•D Service Condition

o Annealed and Cold

Deformed

o Service Condition and

Cold Deformed

100 —

r

V) 400 -it:

- 300 'HL

200

- 100

20 40 60 80 100

PERCENT MARTENSITE

Figure 2. Fracture toughness vs. percent martensite for AISI 304L

and AISI 30^ annealed and AISI 30^ annealed and cold

deformed. ^ 1 1

2

PERCENT MAGNETIZATION, FN

10 20 30 40 50

»

PERCENT MARTENSITE

CO

CO

UJz

I

o

3

O

fid

oc3

H

O

<

a.

Figure 3. Fracture toughness vs. percent martensite for AISI 308L

weld metal.

153

FRACTURE

TOUGHNESS.
Figure

I 1 100 ym

20 ym

SEM photomicrographs of the J-test fracture surface of 304L

specimen with martensite.

154
Figure 5.

1 1 100 ym

SEM photomicrographs of the J-test fracture surface of 308L

specimen in the annealed condition.

155
Figure 6.

I 1

1 00 ym

20 ym

SEM photomicrographs of the J-test fracture surface of 308L

specimen in the service condition.

156
20 ym

Figure 7 . SEM photomicrographs of the J-test fracture surface of 3O8L

specimen in service-and-deformed condition.

157

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CARBON AND NITROGEN EFFECTS ON THE ELASTIC CONSTANTS

OF A STAINLESS STEEL AT 4 K*

H. M, Ledbetter, M. W. Austin, and S. A. Kim

Fracture and Deformation Division,

Institute for Materials Science and Engineering

National Bureau of Standards, Boulder, Colorado

Nine AISI type 30^ stainless steel alloys were studied between

ambient temperature and liquid-helium temperature. The carbonplus-nitrogen contents of these alloys ranged from 0.080 to

0.359 wt./t (from 0.3 to 1.5 at.%)

,

Five elastic constants (the

longitudinal modulus. Young’s modulus, the shear modulus, the bulk

modulus, and Poisson’s ratio) were determined by a pulse echo

ultrasonic method. Except for the bulk modulus, all the elastic

stiffnesses increase slightly with increasing interstitial content.

Existing elastic theory predicts decreases in elastic stiffness

caused by the volume increase. Previously reported ambienttemperature results show such elastic-constant decreases. Thus,

the low temperature results suggest electronic interactions between

interstitial atoms and host-lattice atoms, interactions that

stiffen the lattice more against shear than against dilatation.

1 . INTRODUCTION

The aim of this study was to determine how carbon-and-nitrogen interstitial alloying additions affect the elastic constants of AISI type 304

stainless steel alloys at liquid-helium temperature, 4 K.

Both carbon and nitrogen increase the strength of austenitic Fe-Cr-Ni

alloys. Nitrogen stabilizes the austenite, reducing the tendency toward

martensitic transformations at low temperatures. Carbon functions similarly,

but at higher concentrations it causes sensitization during welding. The

increasing uses of these alloys, together with a dearth of reliable physical

property measurements, prompted the present study.

Because elastic constants are physical properties, they are relatively

insensitive to details of internal structure, which are related to thermalmechanical treatment. However, when determined with sufficient sensitivity,

elastic constants provide important parameters for indicating interatomic

force changes that may occur during alloying. Indeed, small changes in the

elastic constants may indicate large changes in the nature of the interatomic

bonding.

Commercial austenitic stainless steels always contain some carbon and

nitrogen, in amounts up to a few atomic percent. These small foreign atoms

^Intended for publication in Materials Science and Engineering.

159
occupy the octahedral hole of the f.c.c. lattice. This largest f.c.c.

interstice provides a space of radius up to _ 1 )r^ = O.^lr^, where r^

denotes the radius of the host atom on the f.c.c sites. Larger interstitial

atoms must expand the f.c.c. lattice. Both carbon and nitrogen atom sizes

exceed the octahedral hole size: ^c^^^Fe = 0.890/1.271 = 0.70; rj^/rpg = 0 . 855 / 1.271 = 0 . 67 . (The value of the radius of the nitrogen atom is

uncertain. Here, as for carbon, we use the Pauling tetrahedral covalent

radius corrected for coordination number.) The lattice expansion caused by

carbon and nitrogen does not remove the usual metallic properties of

ductility and conductivity; these properties are only reduced. Consistent

with a trend toward the properties of a covalent material, the strength

increases.

Previously, the authors reported a similar study at ambient temperature [1]. They found that carbon and nitrogen decrease all the elastic

stiffnesses: Cj^ = longitudinal modulus, E = Young's modulus, G = shear

modulus, B = bulk modulus, and so on. To explain these results, they invoked

Eshelby's model [2], which predicts elastic softening when lattice defects

increase volume. Carbon and nitrogen increase the volume of these alloys

approximately 0.7% per atomic percent interstitial [3].

Low-temperature studies assume importance for several reasons:

(1) Increasingly, stainless-steel alloys find low-temperature applications;

for example, in magnetic-fusion-energy devices. (2) Usually we consider the

T = 0 K state as a material's ground state. Because this state omits temperature effects, various theoretical calculations and models become simpler.

(3) Carbon-plus-nitrogen strengthening of these alloys depends markedly on temperature [4]. Especially at low temperatures, nitrogen strengthens more

than carbon. The elastic constants enter prominently into most models of

solid-solution strengthening.

2. EXPERIMENT

2.1 Materials and ambient temperature

For brevity, we omit these, which we gave previously [1]. Figure 1 shows a photomicrograph of one of the nine alloys. Although the previous

study used x-ray diffraction mass densities, the present study uses Archimedes's method of mass densities, which are approximately one percent

lower. Also, the carbon-plus-nitrogen chemical compositions in Table 1 differ from those given previously [1]. We consider the current values more nearly correct.

2.2 Low-temperature methods

We achieved cooling to liquid-helium temperature by methods described

previously [5]. At each temperature, approximately at 5-K Intervals, we made

absolute measurements of the velocity of sound in the longitudinal and

transverse directions (v^ and v-^-) in two coolings of the same specimen.

160
RESULTS

Figure 2 shows relative values of the longitudinal modulus

= Pvf (1 )

and the shear modulus

( 2 )

versus temperature, where p is mass density.

Table 1 gives the principal results of the study: the velocities of

sound and the elastic constants for nine AISI type 304 stainless steel alloys

with different carbon-plus-nitrogen contents. Figure 3 shows these results

graphically, based on smoothed sound-velocity results. Table 2 contains the

linear least-squares parameters for the various elastic constants. The last

column in Table 2 gives the percent elastic constant changes per atomic percent of interstitials.

4. DISCUSSION

Figure 2 shows that the elastic stiffnesses soften anomalously during

cooling. This softening, or transition, temperature depends on interstitial

content according to the relationship [6]

where x denotes atomic percent and T^. is in kelvins. As described elsewhere [7],

this transition corresponds to a magnetic transition from the high- temperature

paramagnetic phase to the low-temperature phase that combines antiferromagnetic and spin-glass states, depending on chemical composition and

temperature. Because the antiferromagnetic state involves spin ordering,

cooling through the transition increases volume [8] and decreases elastic

stiffness.

Figure 3 shows the study’s principal results: the 4-K variation of

elastic constants with interstitial content. When compared with the 293“K

results reported previously [1] and shown in Fig. 4, we see a dramatic

difference: the shear-mode elastic constants, the Young’s and shear moduli,

increased with alloying at 4 K; at 293 K they decreased.

Previously, we explained that the ambient-temperature softening by

alloying arose from the volume Increase caused by interstitial carbon and

nitrogen [1]. In that study, we suggested that carbon-iron and nitrogen- iron

electronic interactions might be necessary to understand the shear-modulus

behavior.

Tt = 53.6 = 13.9 x^,N (3)

161
The new low-temperature results underscore the necessity to consider

electronic effects. Elastic models can not explain increases in shear

modulus arising from interstitial atoms that increase volume.

Concerning iron-X (X denotes carbon or nitrogen interstitial atom)

interatomic bonding, we can learn much from previous studies on intermetall ic

compounds. When an X atom occurs in the f.c.c. unit cell, locally the

compound Fe^^x exists. If X atoms cluster, then the compound may achieve the

FeX stoichiometry. In either case, locally there exists an intermetallic

compound. Properties of such compounds include high melting point and high

strength. These properties represent localized-electron (ionic, covalent

bonding) rather than nonlocalized-electron (metallic bonding) systems.

Localized-electron bonding produces certain elastic-constant effects:

one of these is a low Poisson’s ratio. At 4 K, the changes in Poisson’s

ratio due to interstitial alloying effects are approximately five times

larger than those at 293 K. Related to this, we repeat the conclusion of

Kdster and Franz [9]: "Poisson’s ratio depends to a much greater extent on

the conditions of bonding than do the other elastic coefficients."

Another related effect of localized-electron bonding is the stiffening

of shear modulus but not bulk modulus. Diamond represents the paradigm of

localized-electron systems. Although diamond’s shear modulus exceeds iron’s

by a factor of 6. A, diamond’s bulk modulus exceeds iron’s by only 2.6. Thus,

localized-electron bonding corresponds to shear rigidity but not volume

rigidity, to rigid interatomic-bond angles but not rigid lengths.

6. CONCLUSIONS

From this study, there arise several conclusions:

1. Previous studies at ambient temperatures showed that alloying interstitial carbon and nitrogen into Fe-l8Cr-10Ni alloys decreased all

of the usual engineering elastic constants. Except for the sheartype elastic constants, Eshelby’s model explained that these

decreases arise from the volume increase associated with interstitial carbon and nitrogen. The shear-modulus slope (1/x)(AG/G), too

low by a factor of 2, suggested electronic effects.

2. The new M-K results differ dramatically from the 293“K results.

Especially, the shear-mode elastic constants differ. Both Young’s

modulus and shear modulus increase with alloying interstitial carbon

and nitrogen.

3. These new results, coupled with previous ones, suggest strong

electronic interactions between host and interstitial atoms that

affect the elastic constants. These interactions involve localizedelectron (covalent, ionic) bonding that strengthens the alloy

against shear deformation but does not strengthen against volume

deformation.

162
ACKNOWLEDGMENTS

The U.S. DOE Office of Fusion Energy supported this study.

REFERENCES

1. H. M. Ledbetter and M. W. Austin, Mater. Sci. Eng., (1985) 143-1^9.

2 . J. D. Eshelby, J. Appl. Phys., ^ (1954) 255-261.

3. H. M. Ledbetter and M. W. Austin, Dilation of an fee Fe-Cr-Ni alloy by

intersitial carbon and nitrogen, submitted to Metall. Trans.

4. R. P. Reed and N. J. Simon, Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures - VII , NBSIR 84-3000, National Bureau

of Standards, Boulder, Colorado (1984) 13~24.

5. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics, J_4 ( 1 974)

246-249.

6. T. Datta and H. M. Ledbetter, unpublished research.

7. T. Datta, D. Thornberry, E. "R. Jones, and H. M. Ledbetter, Solid State

Commun., ^ (1984) 515-517.

8. V. L. Morruzzi, J. F. Janak, and A. R. Williams, Calculated Electronic

Properties of Metals , Pergamon, New York (1978) 164.

9. W. Koster and H. Franz, Metall. Rev., 6 (1961) 1-55.

163
TABLE 1

Elastic constants at 4 K

C+N

( at . %)

Mass Density

(g/cm3)

^t

( cm/us) (cm/ijs)

G

(GPa) (GPa)

E (GPa)

B (GPa)

Poisson's

Ratio

V

0.34 7.919 0.3214 0.5808 81.80 267.1 209.3 158.0 0.279

0.45 7.909 0.321

8

0.581

1

81 .90 267.1 209.5 157.9 0.278

0.69 7 .900 0.3222 0.5815 82.01 267.1 209.7 157.8 0.278

0.79 7.894 0.3224 0.581

7

82.07 267.1 209.8 157.7 0.278

0.88 7.885 0.3228 0.5820 82.17 267.1 210.0 157.5 0.277

0.95 7.876 0.3232 0.5823 82.27 267.1 210.2 157.4 0.277

1 .29 7.864 0.3237 0.5828 82.40 267.1 210.5 157.2 0.276

1 .38 7.856 0.3241 0.5831 82.49 267.1 210.6 157.1 0.276

1 .47 7.847 0.3244 0.5834 82.59 267.1 210.8 157.1 0.276

TABLE 2

Least-squares parameters for various elastic constants, C, at T = 4 K, where

C = Co + axQ+i^

C Cq (GPa) a (GPa) 100(1/xc+i^)(AC/Co)

B 158.2 - 0.818 -0.52

G 81.6 0.663 0.81

E 208.9 1 .283 0.61

267.0 -0.066 -0.025

V (dimensionless) 0.280 - 0.0025 -0.896

164
C(T)/C(295)

Fig. 1. Photomicrograph of typical studied alloy. Average grain size for

nine alloys equals 88 pm.

Fig. 2. Relative variation of longitudinal and shear moduli for one alloy versus temperature. Elastic softening near 50 K during cooling

results from a paramagnetic to antiferromagnetic/spin-glass

transition.

165
AC/C

Fig. 3- Variation in elastic constants with interstitial carbon-plusnitrogen content for AISI type 304 stainless steel at 4 K.

C+N (at.%)

Fig. 4. Companion graph to Fig. 3 showing ambient-temperature behavior.

166


EFFECTS OF INTERSTITIAL CARBON AND NITROGEN ON THE LOW-TEMPERATURE

MAGNETIC PROPERTIES OF FCC Fe-Cr-Ni ALLOYS

E. R. Jones, Jr., T. Datta, and C. Almasan

Department of Physics and Astronomy

University of South Carolina

Columbia, South Carolina

and

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

We have studied the low- temperature magnetic susceptibility

of a series of fee Fe-Cr-Ni alloys containing varied amounts of

interstitial carbon and nitrogen. Characteristic of antiferromagnets, all specimens displayed a distinct peak in magnetic

susceptibility as a function of temperature. Above the transition temperature, their behavior was well described by a modified

Curie-Weiss law, x = Xq + C/(T + 9). The position of the sus- ceptibility maximum decreased from 47.9 + 0.5 K to 35.0 + 0.5 K

as the C+N content increased from 0.34 to 1.47 at.7o. Similar

behavior was also observed in the transition temperatures determined from the anomaly in the ultrasonic wave velocities.

INTRODUCTION

The low- temperature magnetic properties of austenitic stainless steels

have been of considerable interest since it was learned that they are not

all nonmagnetic as previously thought. Following the observation by

Kondorskii and Sedov^ of antiferromagnetism in a specimen of 18 Cr, 9 Ni

(wt.7o) commercial stainless steel, there have been numerous studies of the

low- temperature structure and properties of the fee austenitic phase (Y ) stainless steels. Ishikawa et al.^ confirmed the long-range antiferromagnetic order at low temperatures in monocrystal fee stainless steel by

neutron diffraction techniques. Evidence for magnetic ordering was also

obtained in Mbssbauer experiments,^’^ in which the single resonance line was seen to broaden as the temperature was lowered through the phase

transition.

167
The additional alloying elements that are found in commercial steels

have significant effects on their magnetic behavior. Extensive studies have

been directed at systematic investigation of the effects of the individual

constituents of the alloys. Warnes and King^ analyzed the data for 30

austenitic stainless steels and special Fe-Cr-Ni alloys whose compositions

fall near the AISI 300 series. A linear relationship was found that predicted the Neel temperature, with an rms deviation of 3.5 K on the basis

of the weight percentage of alloy compositions. The effect of alloying elements in lowering T]^ increases in the order Cr, Ni, Mo, and Si, while the

effect of Mn is to raise T[^, Warnes and King examined alloys whose basic

compositions lie in the range 16-25 Cr , 6-32 Ni, 0-16 Mn, 0-2.4 Si, 0-3.0 Mo,

0-0.14 C, and 0-0.38 N, balance Fe (wt.7o). They reported that the interstitial solutes C+N exert negligible influence on the Neel temperature.^’®

The present study is concerned with the low- temperature magnetic properties of austenitic Fe-Cr-Ni alloys. Particular emphasis is given to the

effects of C+N interstitials on the paramagnetic- to-antiferromagnetic phase

transition, as determined from the peak of the dc susceptibility.

EXPERIMENT

All nine of the specimens studied were 304-type fee stainless steel

alloys. The C+N contents were varied in three steps each while the remaining

constituents were held essentially constant. The detailed chemical compositions are given in Table I. Individual specimens were cut into rectangular

prisms with a typical size of 1.5 mm x 3 mm x 5 mm. The specimens were

cleaned in a solution containing equal parts of HCl and HNO^ plus small

amounts of H 2 SO4 . Ultrasonic agitation was used to speed the cleaning.

Finally, the samples were rinsed in distilled water.

The magnetic measurements were performed with a computer-controlled

superconducting- quantum- interference-device (SQUID) variable- temperature

susceptometer . The specimens were initially cooled to 5 K in zero applied

magnetic field. The susceptibility was then measured over the temperature

range 5 to 200 K in applied magnetic fields of 1, 5, 20, and 50 mT. Thermohysteresis that has been reported in related alloys^ v/as not observed

with any of these samples.

The crystal lattice spacings were measured by X-ray diffraction using

a modified commercial spectrometer. Details of the X-ray measurements

appear elsewhere.

12 13 Sound velocities were determined b)'" a technique detailed elsewhere. ’ Briefly, a pulse-echo method was used to determine the velocity v = 2l/x

,

where t denotes the round-trip time for an ultrasonic pulse between two flat,

parallel specimen faces separated by a thickness, 1. Measurements were made

for both longitudinal and transverse waves.

168
RESULTS

The X-ray measurements indicate that the C+N present in these alloys occur at interstitial sites. These interstitials contribute to a gradual

increase in the average lattice spacing, a^, with increased C+N content.

The linear relationship is obvious in Fig. 1, where we have also included ^ the unit cell. The lattice constant increases at the rate of 7.62 x 10“^ A

per at.7o C+N. A change in lattice spacing of this magnitude would be ex- pected to have a noticeable effect on the magnetic behavior.

The magnetic susceptibility of each of the nine specimens displayed a distinct peak as a function of temperature. This behavior is characteristic

of a paramagnetic- to-antiferromagnetic phase transition.

A

plot of the

reciprocal mass susceptibility versus temperature is shown in Fig. 2. Above

the transition temperature, the susceptibility was well described by a modified Curie-Weiss law;

X(T) = Xq + C/(T + 9) , (1)

where Xq, C, and 9 are constants.

The temperature-independent Xq results from underlying diamagnetism,

Pauli spin paramagnetism, and a saturated ferromagnetic component due

to small ferromagnetic clusters within the alloy. ^ The constants C and 9

represent the Curie 'constant and the paramagnetic Curie temperature. Table II

gives the values of X^, C, and 9 obtained from the computer fit of Eq. 1 for

all nine samples. Figure 2 also shows a plot of l/(x - Xq) versus temperature for sample number 1 in a magnetic field of 20 mT. The line represents

a fit of Eq. 1 to the data. The value of X^ generally decreases as the con- centration of C+N Interstitials increases. This is consistent with expected

changes in the Landau-Peierls and Pauli susceptibilities^'^’^” due to the

lattice dilation and the reduction of the Fermi energy.

The maximum value of the susceptibility is reduced with increased dc

magnetic field. This behavior is similar to that reported by Warnes and

King^ for 20 Cr, 16 Ni (wt,7o) steel. In the present study, however, the

peak was noticeably sharper and, at the fields employed, did not disappear

as happened for the^specimens studied at higher fields by Warnes and King^

and Ishikawa et al. The position of the susceptibility max.imum, which we

take as Tj^, was unaffected by the magnitude of the field over the region

studied, 1 to 50 mT. In contrast, Warnes and King reported a gradual re- duction in T^ with increasing magnetic field for fields from 46 kA/m to

367 kA/m (58 to 460 mT) and a complete disappearance of the peak at still

higher fields.

The positions of the maxima were strongly affected by small changes

in the C+N content, as seen in Fig. 3, where we have plotted normalized

susceptibility, X/Xjj^^j., versus temperature for three of the specimens.

Increasing the C+N content decreased the temperature of the susceptibility

maximum, as shown by the solid circles in Fig. 4. The value of Tj^[ determined from the position of the susceptibility maximum decreased from

169
47.9 ± 0.5 K at 0.34 at.7o C+N to 35.0 ± 0.5 K at 1.47 at.% C4-N. The line

in Fig. 4 corresponds to a least-squares fit of the data, which gives an intercept at 54.0 K and a slope of -12.73 K/at.7o (C4-N).

The transition temperatures were also determined from the anomaly in

1 ft 1 Q the ultrasonic wave velocities. ’ In Fig. 4 we have also plotted the

values of T]^ obtained from the ultrasonic wave velocities. The values of T^

obtained from the longitudinal velocities are larger than the values obtained

from the susceptibility maxima, whereas the values of Tj,^ obtained from the

transverse velocities are slightly lower. The general trend of T^ versus C+N

(at,7o) from the elastic data parallel that of the susceptibility data. A fit

of the ultrasonic data to a linear regression gives slightly different slopes

for the transverse and longitudinal velocities, but within their statistical

uncertainties they agree with the slope obtained from the susceptibilities.

CONCLUSIONS

We have observed a 277o decrease in the magnetic transition temperature

for an increase in C+N content from 0.34 to 1.47 at.7c This considerable

change in the transition temperature may be expected from the change in

lattice parameter. However, this result is in marked contrast to the negligible influence presumed by previous researchers.^ The negative result of

the earlier work may be attributable to the narrow range of variation in

the C+N content in their samples.

The present observations warrant an expanded investigation into the

influence of C+N and other constituents of the Fe-Cr-Ni alloy system. Such

a study is under way.

ACKNOWLEDGMENTS

The work was supported in part by a grant from NSF ISP-80-11451, The

study at NBS was supported by the U.S. DOE Office of Fusion Energy,

REFERENCES

1, E, I, Kondorskii and V, L, Sedov, Sov, Phys . JETP 8, 1104 (1959),

2, Y. Ishikawa, Y, Endoh, and T, Takimoto, J. Phys. Chem, Solids 31, 1225

(1970).

3, Y, Gonser, C, J, Meechan, A, H. Muir, and H. Wiedersich, J, Appl. Phys.

34, 2373 (1963).

4, L. D. Flansburg and N. Hershkowitz, J, Appl, Phys, 41, 4082 (1970),

5, L. A, A, Warnes and H, W. King, Cryogenics 16, 473 (1976),

6, H, W, King and D. C, Larbales tier , Cryogenics 21, 521 (1981),

7, L, A. A, Warnes and H, W, King, Cryogenics 16, 659 (1976),

8, L. A, A, Warnes and H, W. King, Cryogenics 21, 729 (1981),

9, T. Datta, D, Thornberry, E. R. Jones, Jr,, and H, M, Ledbetter, Solid

State Commun, 52, 515 (1984),

170
10 . T. Datta, D, Thornberry, C, Almasan, and E, R, Jones, Jr., Solid State

Commun. ; in press.

11. H. M. Ledbetter and M. W. Austin, Metall. Trans.; in press.

12. H. M. Ledbetter, N. F. Frederick, and M. W. Austin, J. Appl. Phys . 51,

305 (1980).

13. H. M. Ledbetter and M. W. Austin, Mater. Sci. Eng., 70, 143 (1984).

14. S. Chikazumi, Physics of Magnetism (Wiley, New York, 1964), p. 82.

15. C. Kittel, Introduction to Solid State Physics (Wiley, New York, 1976),

5th Ed., pp. 449-452.

16. R. M. White, Quantum Theory of Magnetism (Springer-Verlag, Berlin,

1983), 2nd Ed., pp. 78-96.

17. The curves were fit using the NLIN procedure under SAS R, SAS Institute,

Inc., P.O. Box 8000, Cary, North Carolina 25711.

18. E. W. Collings and H. M. Ledbetter, Phys. Lett. 72A, 59 (1979).

19. H. M. Ledbetter, Res. Mechanica; in press.

171
Table I. Chemical composition (wt.%)

Alloy Cr Ni Mn P S Si Mo Cu C N C+N

1 18.3 10.1 1.57 0.021 0.019 0.50 0.20 0.210 0.038 0.042 0.080

2 0.062 0.043 0.105

3 0.112 0.044 0.156

4 18.0 9.91 1.50 0.022 0.017 0.57 0.205 0.197 0.055 0.137 0.192

5 0.076 0.135 0.211

6 0.096 0.129 0.225

7 20.3 9.97 1.42 0.024 0.016 0.55 0.195 0.200 0.036 0.287 0.323

8 0.062 0.278 0.340

9 0.095 0.264 0.359

Table II. Summary of magnetic measurements

T^ ^0 6^ T

^

-•'N Alloy (K) (10"^mVkg) - (10~^K*m^ /kg) (K) (K)

1 50-139 3.68±0.15 3.4910.19 -13.9611.60 47.9

2 50-199 2.89±0.05 3.6410.09 -13.1510.80 48.1

3 50-200 2.7010.06 3.5010.09 -14.8610.83 47.1

4 50-90 2.3810.10 3.1910.15 -14.7711.52 44.3

5 50-200 2.7110.06 3.6610.10 -13.6110.90 43.8

6 45-144 2.1310.21 4.4710.31 -6.1012.17 42.25

7 40-200 2.9310.03 4.1710.07 -4.2012.47 36.5

8 70-199 2.1710.01 3.0510.05 -7.0110,83 36.4

9 40-100 1.4510.61 4.1310.68 -5.7514.10 35.0

Temperature region for modified Curie-Welss fit,

^Contents obtained from fit of X

= + C/ (T + 0) to the data,

c Ordering temperature obtained from susceptibility maxima.

172
C+N CONTENT (at. %)

Fig. 1. Effect of C+N content on the lattice constant a in fee Fe-Cr-Ni

steel. Insert shows schematic representation o? the unit cell.

Partly covered circles represent corner atoms. Full circles

represent face-centered atoms. Shaded circle represents maximum

diameter of octahedral hole. Dashed circle shows relative size

of C atom in Fe.

173
Fig. 2. The filled circles represent the inverse of the magnetic susceptibility, x» of Fe-Cr-Ni alloy no. 1 as a function of temperature

in a field of 20 mT. A nonlinear behavior is observed above the

transition temperature. The open circles represent the inverse

of X “ Xq’ straight line is the best fit of the data to

X - Xq = C/(T + 6).

174
TEMPERATURE (K)

Fig. 3. Normalized magnetic susceptibility vs. temperature

for three of the Fe-Cr-Ni alloys.

175
)i

: 1 1 1— 0 0.4 0.8 1.2 1.6 C+N CONTENT (at. %)

Fig. 4. Effect of C+N content (at.%) on Neel temperature as determined

by: susceptibility (filled circles); transverse sound velocity

(triangles) ; longitudinal sound velocity (squares)

.

176
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U.S. /JAPAN COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

FIRST DRAFT

FEBRUARY 1986

PROPOSED STANDARD METHOD

FOR TENSILE TESTING OF STRUCTURAL ALLOYS

AT LIQUID HELIUM TEMPERATURE

1 . SCOPE

This document describes the orocedures used in displacement-controlled

tensile testing of structural alloys and weldments at liquid helium temperature, 4 K (-452°F). The format corresponds closely to ASTM standard

designation E 8-83 entitled "Standard Methods of Tensile Testing of Metallic

Materials." The contents include modifications to accommodate the special

needs of cryogenic testing. To measure the tensile mechanical properties

at a temperature of 4 K (±1 K), the specimen is tested at a low strain rate

while fully submerged in liquid helium. This document addresses the need

for the proper cryogenic apparatus and other relevant concerns, such as

specimen size, serrated yielding, adiabatic heating, and strain rate.

2. APPLICABLE DOCUMENTS

ASTM Standards:

E 4 - "Practices of Load Verification of Testing Machines" (ANNUAL BOOK

OF ASTM STANDARDS, vol . 03.01)

E 6 - "Definition of Terms Relating to Methods of Mechanical Testing"

(ANNUAL BOOK OF ASTM STANDARDS, vol. 03.01)

E 8

- "Standard Methods of Tension Testing of Metallic Materials"

(ANNUAL BOOK OF ASTM STANDARDS, vol. 03.01)

177
3. DEFINITIONS

The definitions of terms relating to tension testing in ASTM Standard

E 6 apply to this standard. Two additional definitions are needed:

a. A cryostat , schematically illustrated in Figure 1 , is a chamber

consisting of a loading frame and Dewar assembly for applying loads to test

specimens in low temperature environments.

b. Adiabatic heating is the internal temperature rise of a specimen

owing to tensile testing under conditions such that the heat generated by

plastic work cannot be quickly dissipated to the surrounding cooling medium.

4. SIGNIFICANCE

The significance of tensile tests in general is described in ASTM E 8,

section 4. At 4 K, tensile tests of metallic materials are unique in that

the typical stress-strain curves obtained in displacement-controlled tests

are serrated. The serrations are formed by bursts of unstable plastic flow,

followed by arrests. The plastic deformation occurs locally with sizable

heat evolution at higher than nominal rates of strain, resulting in periodic

reduction of load —or serrated yields.

Note 1

The exact form of the stress-strain curve during unstable

plastic deformation depends on whether load or displacement

control is used. Displacement control should be used for

conventional material property characterization.

Serrated yielding has been observed in a wide variety of metals and

alloys, including austenitic stainless steels, aluminum, and copper alloys,

when test temperatures are 20 K or lower [1-?]. Representative stress-strain

curves for an austenitic stainless steel (Figure 2) show no serrations at

295, 195, and 76 K. Serrations appear only in the extreme cryogenic range

where (a) the volume specific heat tends toward zero, (b) the thermal

178
conductivity rapidly decreases, and (c) the flow strength depends strongly on

temperature.

At 295 K the specimen temperature rise during straining is relatively

small, a few degrees or less at the strain rates normally used. As the

test temperature decreases, the specific heat decreases so that the temperature rise becomes greater for a given amount of work. Moreover, the thermal

conductivity decreases at low temperatures. Thus, when work is done on a

specimen during tensile tests at low temperatures, the resultant internal

heat is less easily dissipated.

These trends become increasingly pronounced at lower temperatures. At

K if tests are conducted at rates normally used at room temperature, the

internal temperature rise is quite substantial. The effects of adiabatic

heating depend on material composition, but in austenitic stainless steels,

for example, temperature rises greater than 50 K have been reported [6], Such

temperature rises may lower the measured yield and ultimate strengths while

increasing the elongation and reduction of area [7]. Therefore, specifications are required for the cooling medium and allowable nominal strain rates

to ensure accurate and reproducible property measurements at 4 K.

Note 2

The most widely accepted explanation for such behavior is

from Basinski [1,2]. His model applies to metals because their

specific heat decreases greatly and their flow strength increases

greatly as their temperature approaches 4 K.

The adiabatic heating mechanism is pictured as follows: Heating during

dislocation movement causes appreciable softening, and deformation becomes

unstable. Low thermal conductivity at 4 K (compared with that at 295 K)

contributes to a localization of the heat to produce load instability. During instability the plastic flow is localized, and the thermal softening due

179
to heat generated by plastic straining is greater than the strain hardening.

Rapid load drops (serrated yields) are associated with each local instability. The instability is eventually arrested at some strain and temperature

where the flow process becomes stable, owing to local specimen strain hardening that distributes the load over an ever-increasing area at a time when

the rise of specific heat precludes continued temperature increases. Once

the flow has become stable again, the specimen cools to ^ K, and with further

deformation, the process repeats.

Localized specimen strain, associated with each load drop, proceeds

along the specimen, in a manner similar to Lueder's band formation at ambient

temperatures

.

5. APPARATUS

5.1 Test Machines

Machines used for tension testing shall conform to the requirements of

ASTM E 4.

At low temperatures the strengths of most metals and alloys increase

significantly, sometimes doubling or tripling between 295 and 4 K. As a

result of such strengthening, higher loads must be applied to the specimen,

grips, load train, and cryostat at cryogenic temperatures, assuming a constant specimen size. If the loads required to determine the yield and tensile strengths fall outside the range of the available testing apparatus (see

ASTM E 4) then the specimen cross-sectional size must be reduced.

5.2 Grips

Any of the gripping mechanisms discussed in ASTM E 8, section 5.2 may be

used for cryogenic tests. The choice depends on specimen type. For the

round-bar and flat specimens, the threaded-end or pin-loaded grips are frequently used at 4 K.

180
Grips may be attached at the bottom of cryostats using spherically

seated bearings, such as those shown in Figures 2 and 3 of ASTM E 8. To

reduce friction at 4 K, the bearing surfaces may be coated with a suitable

dry lubricant, such as solid molybdenum disulfide compounds.

To resist embrittlement at 4 K, the grips should be made of strong,

tough cryogenic structural alloys. Austenitic stainless steels (AISI 304LN),

managing steels (200, 250, and 300 grades), and titanium alloys (Ti-6A1-^V

ELI and Ti-5Al-2.5Sn ELI) have worked well.

5.3 Dimension-Measuring Devices

See ASTM E 8, section 5.3.

5.^ Cryostat

A cryostat capable of retaining liquid helium for the entire duration of

the test must be provided to achieve the desired temperature of 4 K (±1 K).

Cryostat load frames' must be custom-built, but Dewars for A-K service are

available commercially. Factors contributing to efficient cryostat design

are discussed by Fowlkes and Tobler [8]. Several practical designs are described in the literature [8-12], including a rotating-turret cryostat designed for multiple-specimen testing [12].

Either a single- or double-Dewar configuration may be used. In the

double-Dewar arrangement, as shown in Figure 1 , an outer Dewar of liquid

nitrogen is used to provide a jacket for the inner Dewar of liquid helium.

In any case, the Dewar and lines for transferring liquid helium must be vacuum insulated.

5.5 Liquid Level Indicator

A level indicator is required to ensure that the specimen is submerged

in liquid helium at all times during the test. On-off indicators, such as

the carbon-resistor type, may be used to verify that the liquid helium level

181
exceeds some reference point above the specimen. Alternatively, the position

of the liquid level may be continuously monitored with superconducting element sensors of various lengths attached vertically to the cryostat.

5.6 Extensometers

Strain during specimen alignment or tensile testing may be measured

using surface-bonded strain-gage films or detachable extensometers. Information on appropriate apparatus and techniques is given in References 13

through 15.

5.6.1 Strain gages

A strain gage bonded to the specimen in the axial direction may be used

to measure strain during tensile tests at ^1 K. Strain-gage films are produced in great variety, but not every kind is suitable for i|-K service. An

appropriate gage active element, backing material, and bonding agent should

be selected on the basis of literature precedents and manufacturers’ recommendations. One approach that has proven quite successful is a nickelchromium alloy having a self-temperature-compensated active element and a

glass-fiber-reinforced epoxy-phenolic resin with a fully encapsulated matrix

(350 q) [14].

To compensate for any misalignment, two gages may be used, aligned in

the axial direction on diametrically opposite specimen surfaces. Each gage

should be wired to a dummy bridge using the three-wire temperature-compensating

hookup as shown in Figure 3. A low excitation voltage of about 1.25 V is

preferred for 350-Q gages at 4 K.

Calibration of strain gages at room temperature requires calibration of

the electronics to changes in resistance. Calibration of strain gages at 4 K

requires consideration of the temperature dependence of the resistance in the

182
electronic calibration. Typical full-scale calibration ranges are Mo at room

temperature and 2% at ^ K.

5.6.2 Detachable Strain-Gage Extensometers

With careful materials selection and design, demountable extensometers

can be built to resist thermal shock and moisture condensation for operation

at ^ K. Unlike directly bonded strain gages, these extensometers can be

reused [14-15].

The most common extensometer of this type is the clip-on gage, which may

be custom-built or purchased. A satisfactory example is the cantilever beam

gage [15]. The extension of the specimen within its gage length is sensed by

the cantilever beam gage, which is clipped to the specimen between retaining

pins that are fixed to the specimen. Such extensometers use four strain-gage

films bonded to the gage in a Wheatstone bridge arrangement.

The extensometer must be calibrated directly at room temperature and at

4 K using a suitable calibration device. For calibrations at 4 K, a barrel

micrometer fitted with vertical extension tubes can be used with clip gages

mounted at the bottom end and immersed in liquid helium.

Extensometers that use capacitance measurement to monitor strain have

also been used. The overlapping cylinder type has an extended strain range,

capacitance output that is linear with strain, and adjustable sensitivity.

The parallel-plate type has high sensitivity, but its output must be compensated for the hyperbolic dependence of capacitance on strain.

Note 3

Once the calibration is known and proven to be accurate,

linear, and reproducible, then room temperature checks can be

performed prior to each test series as an indirect verification of the 4-K calibration. However, the calibration at 4 K must be directly confirmed periodically and whenever damage

is suspected, repair has been performed, or gages replaced.

183
6. TEST SPECIMENS

6.1 General

Any specimen type cited in ASTM E 8 may be used for tests at 4 K. The

appropriate specifications for specimens (e.g., dimensions, gage length,

location, surface finish) are given in ASTM E 8, sections 6.1 through 6.17.

Specimens taken from sheet or wire products having relatively small

cross-sectional areas can be tested in typical cryogenic apparatus. However,

specimens from relatively thick plate or bar products may need to be machined

to a reduced cross-sectional area so that the maximum load capability of the

cryogenic test apparatus is not exceeded.

6.2 Round-Bar Specimens

6.2.1 Standard specimen

The recommended standard specimen for 4-K tensile testing is the 6.25-mmdiameter bar specimen. A specimen of this type with threaded ends and a gage

length of 25.4 mm is shown in Figure 4. Its proportions are similar to those

of the standard 1 2.5-mm-diameter specimen of ASTM E 8. A schematic with

relevant notes and specifications on tolerances is given in Figure 8 of ASTM

E 8-83.

Note 4

The 6.25-mm-diameter specimen replaces the 1 2 . 5-mm-diameter

specimen of ASTM E 8 section 6.2 as standard because of its size.

In tests of strong alloys at 4 K, the 1 2. 5-mm-diameter specimen

requires high loads. For example, about 210 kN is required to

test AISI 304LN (ultimate strength = 1690 MPa at 4 K), whereas

100 kN is the upper limit for most cryogenic test machines.

6.2.2 Subsize specimens

If it is not appropriate to use the standard 6.25-mm-diameter specimen,

then another size or geometry may be selected, following ASTM E 8-83 guidelines. For any subsize specimen: (a) the gage length for measurement of

184
elongation should be four times the specimen diameter and (b) the average

grain size of the test material should be reported.

7. PROCEDURES

7.1 Specimen Installation

The specimen must be installed in the gripping fixture such that the instrumentation wires are free of tension and have enough slack so that they

will be neither stretched nor crimped during cryostat assembly and testing.

7.2 Alignment

Proper alignment is essential to avoid unwanted bending of the specimen.

The longitudinal axis of the specimen should coincide with the centerline of

the test machine heads. Cryostats may be designed with adjustable loading

columns or posts to facilitate the overall alignment.

Alignment may be assessed at room temperature with specimens instrumented

with several axial strain gages or extensometers. Axial strain readings from

any point around the circumference of the specimen in' the gage length area

must agree within 5/5 of the average value of all readings. During the

process of alignment, the applied loads should be kept below one-third of the

proportional limit of the test material.

7.3 Cool-down Procedure

a. Remove any moisture present on the apparatus prior to cool-down by

drying thoroughly with a pressurized air jet or heat gun. Ice formations may

block cryogenic transfer lines or cause erratic loading behavior if buildup

occurs between various parts of the specimen, grips, and load train.

b. Assemble the Dewar and precool the apparatus by transferring liquid

nitrogen into the cryostat. After equilibrium is reached at 77 K, be sure to

remove all of the liquid nitrogen (or transfer it to the outer Dewar if a

double-Dewar arrangement is used).

185
c. Transfer liquid helium into the cryostat until the specimen is fully

submerged. Testing may begin after the system has reached equilibrium at

4 K. The specimen must be covered with liquid helium at all times during the

test to ensure that the temperature remains constant (±1 K) and free of gradients along its entire length.

Note 5

The heat transfer characteristics of gases are poorer than

those of liquids. The thermal conductance of the specimen and

the heat capacities of both the specimen and gas change with temperature. Therefore, a liquid helium environment is required to

minimize internal heating effects.

7.4 Speed of Testing

The speed of testing at 4 K must be specified in terms of the nominal

rate of straining. Any convenient speed may be used up to a stress of onehalf the yield strength; thereafter, the rate of straining must not exceed

5x10"^ s"’’ .

Some test machines are equipped with pacing or indicating devices for

measurement and control of strain rates. In the absence of such a device,

the average rate of straining can be determined by observing the time required to effect a known increment of strain.

Note 6

Rates of strain exceeding 5 x s~^ may cause significant

temperature control problems and inaccurate property measurements.

Tests are currently in progress to determine the effects more precisely. The present recommendation is based on the current practice of NBS. It may be revised when additional strain-rate

sensitivity data become available.

7.5 Measurement of Mechanical Properties

a. Measure the yield strength of the specimen according to the Q.2%

offset method, following ASTM E 8, section 7.3.1. If a load drop occurs

before the point of intersection of the stress-strain curve and the 0.255

186
offset line, then report the stress at the highest load preceeding the inter

section point as the yield strength.

b. Use a strain-measurement device (such as a strain gage; strain-gage

clamp-on extensometer ; linear variable differential transformer ; or capacitance sensor) to obtain yield strength data at 4 K.

Note 7

Normal efficient cryostat and load trains for testing

at ^ K are comparatively "soft." The effective modulus of the

load-train system is low and depends on the liquid helium level

and the degree of temperature stabilization that is achievable

throughout the load system in the temperature range 295 to ^ K.

Therefore, attempts to measure yield strength from offsets of

load versus time have large potential error. This technique is

not recommended.

c. Measure the tensile strength, elongation, and reduction of area

according to ASTM E 8, sections 7.5 to 7.7, respectively.

d. Round the calculated test data according to ASTM E 8, section 7.8.

e. If necessary, discard data and test replacement specimens as per

ASTM E 8, section 7.9.

8. TEST REPORT

8.1 Individual Specimens

The test report shall include the following data for each individual

specimen

:

a.

b.

working

c

.

d.

e.

Test temperature and environment

Specimen geometry, size, and orientation relative to

directions of the original stock

Nominal rate of straining

Yield and ultimate strengths

Elongation and reduction of area

the principal

187
f. Optional data, such as the original and final specimen dimensions

(areas and gage lengths), the measured value of Young’s modulus, and the

material's average grain size.

8.2 Replicate Specimens

If replicate specimens are tested, the report should include the number

of replicate specimens tested and the average values of all mechanical property measurements with a measure of scatter.

8.3 Substandard specimens

For subsize specimens, the average grain size of the material must be

given.

8.4 Anomalous Behaviors

Any anomalies observed with respect to test procedures, specimen behavior, or failure location must be described.

9. PRECISION AND ACCURACY

See ASTM E 8, section 8.

10. REFERENCES

1. Basinski, Z. S., "The Instability of Plastic Flow of Metals at Very Low

Temperatures", Proceedings of the Royal Society , vol. A 240, 1957, pp.

229-242.

2. Basinski, Z. S., Australian Journal of Physics , vol. 13, I960, pp.

354-358.

3. Powell, G. W. , Marshall, E. R., and Backofen, W. A., Transactions of

the ASM , vol. 50, 1958, pp. 478-497.

4. Wessel, E. T., Transactions of the ASM , vol. 49, 1957, pp. 149-172.

5. Read, D. T. and Reed, R. P., "Heating Effects During Tensile Tests of

AISI 304L Stainless Steel at 4 K," Advances in Cryogenic Engineering —

Materials , vol. 26, Plenum, New York, 1980, pp. 91-101.

6. Ogata, T., Ishikawa, K., Hiraga, K., Nagai, K., and Yuri, T., "Temperature Rise During the Tensile Test in Superfluid Helium," Cryogenics , vol. 25, no. 8, August 1985, pp. 444-446.

188
7. Dobson, W. G. and Johnson, D. L., "Effect of Strain Rate on Measured

Mechanical Properties of Stainless Steel at ^ K," Advances in Cryogenic

Engineer ing- -Mater ials , vol. 30, Plenum, New York, 1984, pp. 185-192.

8. Fowlkes, C. F. and Tobler, R. L., "Fracture Testing and Results for a Ti-6A1-4V Alloy at Liquid Helium Temperature," Engineering Fracture Me - chanics , vol. 8, 1976, pp. 487 - 500 .

9. Reed, R. P., "A Cryostat for Tensile Tests in the Temperature Range 300°

to 4°K," Advances in Cryogenic Engineering , vol. 7, Plenum, New York,

1962, pp. 448-454.

10. Hartwig, G. and Wuchner, F., "Low Temperature Mechanical Testing Machine," Review of Scientific Instruments , vol. 46, No. 4, 1975, pp.

481-485.

11. Novikov, N. V., "Mechanical Property Measurement Techniques of Structural Materials at Cryogenic Temperatures," Advances in Cryogenic Engi - neering , vol. 22, Plenum, New York, 1977, pp. 113“118.

12. Horiuchi, T., Shimada, M., Fukutsuka, T., and Tokuda, S., "Design and

Construction of an Apparatus for Testing Materials at Cryogenic Temperatures," Proceedings of the Fifth International Cryogenic Engineering

Conference, IPC Science and Technology Press, Guildford, Surrey, England, 1977, pp. 465-468.

13 . Sparks, L. L., "Temperature, Strain, and Magnetic Field Measurements,"

Chapter 14 in: Materials at Low Temperatures , R. P. Reed and A. F.

Clark, eds., American Society for Metals, Metals Park, Ohio, 1983, pp.

555-559.

14. Tobler, R. L. and Shepic, J. A., "Design and Performance of a Ring-Shaped Clip Gage," Journal of Testing and Evaluation , vol. 13, No.

4, 1985, pp. 299 - 302 .

15 . Read, D. T. and Tobler, R. L., "Mechanical Property Measurement at Low

Temperatures," in Advances in Cryogenic Engineering—Materials , vol. 28,

Plenum, New York, 1982, pp. 17-28.

189
Applied Load

Figure 1. A cryostat with double-Dewar arrangement for tensile testing at 4 K.

190
TENSILE

STRESS

(GPa)

Figure 2. Engineering stress-strain curves for an Fe-l8Cr-10Ni

austenitic stainless steel at 295, 195, 76, and 4 K.

191

ENSILE

STRESS

(ksi)
iK 350-Q resistors are used in dummy legs

Figure 3. Method of strain gage wiring for cryogenic tests.

192
Figure H

12.7 I 12 . 6

^

6.30

6.25*"

0.41

REDUCED

SECTION

38.1

71.4

Dimensions are in millimeters

A 6 . 25-mm-diameter tensile test specimen with threaded ends

(see Fig. 8 of ASTM E 8-83 and all relevant notes).

193
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U.S. /JAPAN COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

SECOND DRAFT

February 1986

PROPOSED STANDARD METHOD

FOR FRACTURE TOUGHNESS TESTING OF STRUCTURAL ALLOYS

AT LIQUID HELIUM TEMPERATURE

1 . SCOPE

This method covers the equipment, procedures, and data analysis used for

determining the fracture toughness of structural alloys and weldments at

liquid helium temperature, K (-452°F). The procedures described here

represent modifications of the Kj^ and plane strain fracture toughness

tests as described in ASTM Standards E 399-83 and E 8l3“8l, to accommodate

the special requirements of cryogenic testing. Since materials that show

ductile tearing in areas of normal thickness are usually used for liquid

helium applications, the Jj^ test procedure is given primary emphasis in this

method.

2. APPLICABLE DOCUMENTS

See ASTM E 399-83, section 2 and ASTM E 813-81, section 2.

3. SUMMARY OF METHOD

See ASTM E 813-81, section 3*

To achieve the desired temperature of 4 K, the specimen is tested

while fully submerged in liquid helium using a cryostat (described in 6.1).

The pin-loaded compact specimen is preferred, because it offers the largest

fracture toughness measurement capacity for its size of any of the common

fracture specimens. A relatively low load applied to the compact specimen

195
produces a large stress intensity factor. This allows a minimum mass to be

used for grips, cryostat frame, and loading train, which significantly

reduces the liquid helium consumption and cost per test.

Note 1

Three-point bend specimens may also be used, but they require

more mass for an equivalent ligament size. The longer span of

the bend specimen transverse to the load axis necessitates more cumbersome loading fixtures, which may be disadvantageous for some cryostats.

4. SIGNIFICANCE

See ASTM E 813-81, section A.

The type of load-deflection behavior exhibited by a material determines

whether Jj^ or Kj^ can be directly measured, as indicated in Figure 1. The

parameter 'Jlc is used to characterize materials where the load-deflection

behavior is nonlinear (i.e., elastic-plastic or fully plastic) and where the

plastic zone at the crack tip is relatively large compared with the crack

size and specimen dimensions. The parameter Kic is used to characterize

materials where the load-deflection behavior is linear or nearly so and

where the plastic zone at the crack tip is small compared with the crack size

and specimen dimensions.

The two parameters, Jj^ and Kj^, are interconvertible under appropriate

circumstances. Together, they are suitable for characterizing a broad range

of materials. The values are routinely converted to obtain estimates of

via the relation: Kj^ = (see ASTM E 813-81, section 9.4). However,

for many austenitic alloys the occurrence of linear-elastic behavior has not

been demonstrated and could be expected only in extremely thick sections, if

196
at all. It must be appreciated that estimations of Kic from Jjq are often

performed merely for the purposes of comparison, and this does not necessarily imply that the given material will display a linear-elastic fracture.

The fracture behavior of ductile alloys at 4 K may be affected by

adiabatic heating during rapid loading [1-^]. Therefore, for ductile

alloys is measured at low loading rates (see section 8.2) to minimize any

adiabatic heating.

A number of alloys of intermediate toughness can not be tested

satisfactorily according to the E 399-83 and E 813~81 methods. In these

materials, fast fracture occurs prior to the full development of the required

four-point resistance curve, but the load-deflection behavior is sufficiently

nonlinear to invalidate a test analysis. In such cases, increasing the

specimen thickness in subsequent tests may result in a vali d Kic test.

There is, however, no assurance that valid results for practical specimen

sizes can be obtained by the procedures of this method.

5. TERMINOLOGY

See ASTM E 616-82, ASTM E 8l3"8l, section 5, and ASTM E 399-83, section 5.

6. APPARATUS

See ASTM E 8l3~8l, section 6.

6.1 Cryostats

A cryostat must be provided to cool and test the specimen at 4 K. A

cryostat consists of a load frame and at least one vacuum-insulated Dewar for

containing liquid helium. Cryostat designs vary according to function and

ingenuity, but the goal is always to achieve a mechanically and thermally

efficient design. An analysis has been given of factors entering into

optimum cryostat design [5], and several satisfactory designs are reported in

197
the literature [5-8]. Single- or double-Dewar arrangements may be used. In a

double-Dewar arrangement, the vacuum-insulated inner Dewar contains liquid

helium while the outer Dewar contains liquid nitrogen, as shown in Figure 2.

6.2 Grips

Flat-bottomed clevis grips conforming to the design shown in Figure 3 of

ASTM E 813-81 should be used for compact specimens. The flats must be free

of indentations to avoid undesirable effects of , friction on the test results.

High-strength materials with adequate fatigue resistance and toughness must

be used. Managing steels (i.e., 200, 250, and 300 grades) and AISI 304LN

have given satisfactory service at ^1 K [5,7].

6.3 Clip Gages

See ASTM E 8 I 3

- 8 I, sections b.^l.l and 6.4.2, and Reference 9.

The clip gage must be able to sustain repeated thermal shock, moisture

condensation, and mechanical fatigue while rendering, accurate deflection

measurements at 4 K. This is accomplished by careful selection of gage

materials and fabrication techniques [9]. Commercially available cryogenic

strain gage films are used in a typical Wheatstone bridge arrangement. The

gage accuracy and linearity must first be determined by direct physical

calibrations at room temperature and at 4 K. This can be done using a barrel

micrometer fitted with vertical extension tubes with the gage mounted at one

end and immersed in liquid helium for calibration at 4 K [9].

Prior to each cryogenic test series, the gage calibration must be

checked. Checks may be performed at room temperature, as an indirect

verification of the 4-K calibration. If so, the 4-K calibration should be

verified periodically, or whenever clip gage damage is suspected.

198
SPECIMEN CONFIGURATION, DIMENSIONS, AND PREPARATION

7.1 Specimen Size

See ASTM E 813 - 8 I, section 7.1

7.2 Three-Point Bend Specimen

See ASTM E 8l3f section 7.2

7.3 Compact Specimen

See ASTM E 81 3“8l , section 7.3

7 .^ Fatigue Crack Length Compatibility

For a ^Ic test the initial crack-length-to-width (a/W) ratio must fall

in the range 0.5 to 0.75; 0.6 is thought to be optimum. For some tests it

may be desirable to choose a crack length in the range 0.5 ^ a/W $ 0.55.

This maintains compatibility with ASTM E 399-83 and allows a Kj^ analysis to

be employed in case of unexpected low-toughness behavior. For 25 . 4-mm- thick

specimens (W = 50.8 mm) of austenitic stainless steels, the notch should be

machined to a depth about 4 mm shorter than the intended final crack length,

because crack front curvature typically amounts to 3 mm at 4 K. The

intention is to minimize the time consumed in fatigue cracking and to

anticipate the occurrence of crack curvature while finishing in the desired

a/W range.

Note 2

The clip gage may be attached between razor blades spotwelded to the specimens or between integral knife edges

machined into the specimen.

7.5 Alternative Specimens

See ASTM E 813, section 7.5.

Experience indicates that a 25.4-mm thickness (W/B = 2) is sufficient for

most of alloys tested at 4 K (e.g., Ti-5Al-2.5Sn and T1-6A1-4V, all ferritic

199
steels, and austenitic stainless steels with yield strengths greater than

600 MPa).

Side grooves may be employed in attempting to improve the crack

front shape, to lower the J-Aa regression line slope, or to improve the

agreement between the true physical crack extension and the calculated value

from compliance unloading tests. The rules for side grooving are stated in

ASTM E 813~81, section 7.5.1. However, if side grooves are used, the Kj^

analysis procedures of ASTM E 399-83 cannot be employed, even if a linear

elastic fracture is observed. In such case, a retest may be conducted

using specimens with plane sides.

7.6 Fatigue Precracking

All specimens shall be fatigue precracked to establish a sharp

crack front. The maximum stress intensity factor during the fatigue crack

extension shall not exceed:

1 /P

a. 0.005 mm times Young’s modulus, E, or

b. 60^ of Kj^ (assuming linear-elastic behavior)

Requirement (a) is established by ASTM E 813 - 8 I, sections 7.6.1 and 7.6.2;

requirement (b), by ASTM E 399-83, section A 2.11.2.

The load range (AP) shall not be less than 90^ of the maximum load.

Fatigue precracking may be done at a temperature different from the

test temperature, providing that: (a) the maximum fatigue loads are reduced

by the ratio of the yield strength at the fatigue temperature to that at the

test temperature, and (b) there is no martensitic phase transformation of the

material at these temperatures . For materials exhibiting martensitic

phase transformations, the fatigue precracking can be done at a temperature

200
different from the test temperature only if it can be demonstrated that this

does not affect the toughness test result.

Note 3

Sizable reductions of the maximum load during fatigue cracking are to be avoided because they may retard the crack growth

rate and increase the degree of crack front curvature.

For heat treating, minimum crack extension from the notch, and notch

radius, see ASTM E 8l3~8l, sections 7.6.3» 7.6.M, and 7.6.5.

8. COOLING AND TESTING PROCEDURES

8.1 Cooling Procedure

The cryostat must be capable of achieving and maintaining a uniform

specimen temperature of 4 ± 1 K for the duration of the test. This is

accomplished by completely submerging the specimen and clip gage in liquid

helium and monitoring the liquid helium level with a suitable meter [1,2].

Sufficient time at temperature must be allowed prior to testing to be sure

that the thermal equilibrium has been reached.

Note 4

The specimen may be precooled to 77 K using liquid nitrogen,

prior to cooling to 4 K with liquid helium. Various cool-down

procedures are discussed in the literature [5-8]. One must carefully and properly purge the transfer line and remove any moisture from the cryostat to prevent the formation of ice blocks.

8.2 Mechanical Test Procedure

Specimens shall be loaded at a rate such that the time taken to

reach the load at first deviation from the linear portion of the record lies

between 0.5 and 10 min. For other details of test procedure, see ASTM

E 8l 3“81 , section 8

.

201
I

NOTE 5

Because fracture testing in liquid helium environments is

time-consuming and costly, a single specimen test technique is

recommended. The most familiar technique to date is the singlespecimen unloading compliance test technique. See ASTM E 813,

section 8.3.4 and Annex A1 and the Appendix to this document.

9. CALCULATION AND INTERPRETATION OF RESULTS

Depending upon the load- versus-deflection behavior, either the

(ASTM E 813~81) or the Kj^ (ASTM E 399-83) analysis procedures should be

employed.

For specimens exhibiting highly nonlinear load-versus-deflection curves,

the test procedure based on the resistance curve technique should be

used. See ASTM E 8l3~8l, section 9.

Note 6

The load-deflection curves of materials tested at 4 K are unique in that, they typically exhibit serrations. These serrations may represent either brittle fracture instabilities or ductile flow instabilities due to adiabatic heating [10]. Calculate the area under the load-deflection curve as given by the

autographic recorder in the usual manner. In resistance curve determinations for ductile alloys of medium or high toughness,

small errors are introduced into the resulting J-integral values, owing to the probability that the X-Y recorder can not accurately follow the instability event. However, the errors are considered negligible. If the serrations are due to brittle

fracture instabilities (as verified by scanning electron microscopic examination of failure mode), a valid resistance curve can

not be obtained by the J-integral methodology.

For specimens exhibiting linear-elastic or nearly linear-elastic loading

behavior, the test technique should be used. The notch cutout modifications for load-line displacement measurements instead of the usual edge

displacement measurements do not invalidate the Kj^ calculations at 4 K

according to this method. See ASTM E 399-83, section 9.

10. REPORTING

See ASTM E 8l3~8l, section 10 and ASTM E 399-83, section 10.

202
n. PRECISION AND ACCURACY

See ASTM E 8l 3-81 , section 11 and ASTM E 399-83, section 11.

1?. REFERENCES

1. Ogata, T., Hiraga, K., Nagai, K., and Ishikawa, K., "A Simplified Method

for Charpy Impact Testing near Liquid Helium Temperature," Cryogenics , vol . 22, no. 9, September 1982, pp. 481-482.

2. Dobson, W. G. and Johnson, D. L., "Effect of Strain Rate on Measured

Mechanical Properties of Stainless Steel at 4 K," Advances in Cryogenic

Engineering—Materials , vol. 30, Plenum, New York, 1984, pp. 185-192.

3. Read, D. T., and Reed, R. P., "Heating Effects During Tensile Tests of

AISI 304L Stainless Steel at 4 K," Advances in Cryogenic Engineering

Materials , vol. 26, Plenum, New York, 1980, pp. 91-101.

4. Ogata, T., Ishikawa, K., Hiraga, K., Nagai, K., and Yuri, T., "Temperature Rise During the Tensile Test in Superfluid Helium," Cryogenics , vol.

25, no. 8, August 1985, pp. 444-446.

5. Fowlkes, C. F. and Tobler, R. L., "Fracture Testing and Results for a T1-6A1-4V Alloy at Liquid Helium Temperature," Engineering Fracture

Mechanics , vol. 8, 1976, pp. 487-500.

6. Read, D. T. and Tobler, R. L., "Mechanical Property Measurements at Low

Temperatures," Advances in Cryogenic Engineering--Materials , vol. 28,

Plenum, New York, 1982, pp. 17-28.

7. 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, no. 11, November 1983, pp. 603-607.

8. Takeuchi, M., Shoji, T., Takahashi, h., and Anayama, T., "Evaluation of

Fracture toughness of Cryogenic Structural Materials," Report of the

Research Institute for Strength and Fracture of Materials , Tohoku

University, Sendai, Japan, vol. 17, nos. 1 and 2, March 1984, pp. 1-12.

9. Tobler, R. L. and Shepic, J.A., "Design and Performance of a Ring-Shaped

Clip Gage for Fracture Mechanics Testing," Journal of Testing and

Evaluation , vol. 13, no. 4, July 1985, pp. 299-302.

10. Tobler, R. l., "Ductile Fracture with Serrations in AISI 310 S Stainless

Steel at Liquid Helium Temperature," Elastic-Plastic Fracture: Second

Symposium, Volume 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. 11-763,

11 - 776 .

203
LOAD,

P

LOAD,

P

•

LOAD,

DISPLACEMENT, 8

(a)Linear-elastic

DISPLACEMENTS

(b)Elastic-plastic (failure before limit load)

(c)Fully plastic (exhibits a limit load)

Figure 1. Load-deflection behavior observed in fracture toughness tests,

204
1 - Dewar pivot frame 19 - Isolator compression ring

2 - Specimen clevis 20 - Adapter bar

3 - Shear pin 21 - Test machine frame

4 - Pivot plate 22 - Load cell

5 - Outer Dewar top 23 - Clip gage

6 - Inner Dewar clamp 24 - Stainless steel Dewar

7 - Transfer hose adapter 25 - Stainless steel Dewar

8 - Adapter plate 26 - J-integral compact tension

9 - Column adapter fracture specimen

10 - Compression column 27 - Seal ring

11 - Clevis pull rod 28 - Liquid helium level meter

12 - Pull rod keeper 29 - Socket head cap screw

13 - Pull rod split ring isolator 30 - Socket head cap screw

14 - Load plate 31 - Socket head cap screw

15 - Pull rod coupling 32 - Socket head cap screw

16 - Load cell adapter 33 - Nut

17 - Load column 34 - Threaded rod

18 - Load cell adapter 35 - Shear pin

Figure 2. Cryogenic fracture toughness test facility with the double-Dewar

arrangement [5]

.

205
APPENDIX

THE USE OF UNLOADING COMPLIANCE TO DETERMINE RESISTANCE CURVE BEHAVIOR

See ASTM E 8l3~8l, sections A1.1 through A1.3.

The single-specimen unloading compliance technique has been successfully

applied for tests at 4 K [A1]. However, because of the existence of

serrations in the load deflection curves of cryogenic alloys at 4 K, the

number of J-Aa points on the resistance curve may be limited. Partial

unloadings should be performed at points of arrest just after each unstable

load drop, as illustrated by the example in Figure 1A.

Comparison of the measured physical crack lengths and inferred

crack extensions from the unloading compliance technique will usually

reveal a disagreement that increases significantly as the test proceeds.

This error may persist even with the use of side grooVes. Since this error

diminishes at lower Aa values, the value obtained by extrapolation of the

unloading-compliance R-curve is still deemed valid by this method [A2]

.

However, unless the agreement between the amount of crack extension

predicted by the unloading compliance method and the average measured value

is within 155S, the resistance curve slope is considered invalid and should

not be reported.

REFERENCES

A1 . Shimamoto, S., Nakajima, H., Yoshida, K., and Tada, E., "Requirements for

Cryogenic Structural Alloys for Superconducting Magnet Cases," Advances

in Cryogenic Engineering—Materials , vol. 32, Plenum, New York, 1986, to

be published.

A2. Tobler, R. L. and Shepic, J.A., "Design and Performance of a Ring-Shaped

Clip Gage for Fracture Mechanics Testing," Journal of Testing and

Evaluation, vol. 13, no. 4, July 1985, pp. 299-302.

206
Deflection, J, mm

Figure lA. Typical compact specimen loading behavior at 4 K, showing

serrations and points of unloading compliance method.

207

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Minutes of

LOW TEMPERATURE STRUCTURAL MATERIALS AND STANDARDS WORKSHOP

Tohoku University, Sendai, Japan

March 7, 1986

The Workshop on Low Temperature Structural Materials and

Standards was held March 7, 1986 at Tohoku University. Hosting

this meeting was the Research Institute for Strength and

Fracture of Materials (RISFM) . Three researchers from the US

participated, so the workshop was conducted in English. The

program consisted of: 1) opening remarks and a business session

on fracture test standardization, 2) a seminar entitled Recent

Topics in Deformation and Fracture of Structural Materials at

Cryogenic Temperatures, and 3) closing remarks and conclusion

of business.

The opening remarks were made by Dr. Takahashi (RISFM) . Plans to provide more scientific support for a fracture

toughness testing standard at 4 K were then discussed by

Mr. Tobler (NBS) . At the last meeting (January, 1986) a test

matrix was proposed to experimentally resolve ten technical

items relating to various aspects of standardization. The

original matrix called for a total of 88 tests to be conducted

using various austenitic stainless steels. But the desired

program was too large to be executed by one organization because

of limited time and test facilities. Therefore a more general

cooperative effort is necessary. At this meeting some modifications to simplify the original proposal were described

and assistance of a number of research laboratories was requested

to complete the testing by the end of 1986. The attendees were

invited to consider their decisions to participate and to respond

at the end of the workshop.

The seminar consisted of 13 technical presentations with

preprints available in most cases. The preprints or copies of

the speaker’s viewgraphs were collected for distribution later

as a record of the proceedings.

The final session on the business of fracture testing

for standard development was chaired by Dr. Shoji (RISFM). After

some discussion, specific tests on various materials were

delegated to be carried out in cooperation with RISFM where

Mr. Tobler is presently a visiting exchange scientist.

209
Responsibility for testing the 25Cr - 15Ni steel naturally passed

to JAERI (Mr. H. Nakaj ima) as the only laboratory capable of

testing the large 3T compact specimens. Kobe Steel (Dr.

M. Shimada) then agreed to test the AISl 304 type steel in view

of some mutual interest in the material. Testing of the AISI

316 LN is still unresolved, but might be done at NRIM (Dr. Ogata)

if machine time is available. The revised plan for this program

is summarized in the attached matrix which shows 69 tests to be

conducted by the cooperating laboratories. All of the test

materials have been obtained, except for the AISI 304 stainless

steel which will be delivered to us at the end of May.

REVISED EXPERIMENTAL TEST MATRIX

LABORATORY MATERIAL TYPE OF SPECIMEN

Compact Specimens (items 1 to 8)

JAERI Fe-25Cr-15Ni

0.5 T IT 1.5 T 3 T

NRIM AISI 316LN

48-4

4 8 4 -

Kobe Steel AISI 304 4 8 - -

JAERI and AISI 316LN

Bend Specimens (25 x 25 x 150 mm; item 6)

5 RISFM AISI 304 5

RISFM AISI 310S

Compact Specimens (items 9 and 10)

1 T

3

AISI 316LN 6

Fe-18Cr-3Ni-13Mn 6

NOTE: Total tests = 69

210
WELDING

211

WELVING PROGRAM

LEADER; T. A. Siewert

CONTRIBUTORS: C. N. McCowan, H. I. McHenry, P. T. Purtscher, R. P. Reed,

R. L. Tobler, D. P. Vigliotti

OBJECTIVES

• Investigation of the metallurgical factors that affect the mechanical

properties of stainless-steel weldments at cryogenic temperatures

• Contributions to the development of improved filler metals for welding

stainless steels for use at 4 K

• Evaluation of the mechanical properties of weldments at M K

• Development of methods of detecting defects and evaluation of their

significance in stainless-steel weldments

RESEARCH PAST YEAR (1985)

• Clarification of the effects of manganese (1.5 to 10 wt./5), nitrogen

(0.04 to 0.26 wt.^), and molybdenum (0 to 4 wt.^) on SMA weld properties

on the basis of AWI 308-type welds

• Determination of the 4-K yield strength and fracture toughness of an

Fe-1 8Cr-20Ni-5Mn-0 . 1 6N fully austenitic weld

• Determination of the 4-K yield strength, fracture toughness, and fatigue

properties of an Fe-1 8Cr-1 6Ni-6 . 5Mn-2. 4Mo fully austenitic weld

• Comparison of the strength-toughness relationships for welds evaluated at

NBS over the past decade with corresponding base material properties

RESEARCH THIS YEAR (1986)

• Comparison of the 4-K yield strength and fracture toughness of electron

beam and laser welds in AISI 3l6LN-type base material

• Development of a model to relate the 4-K weld fracture toughness to the

microstructure

• Determination of the 4-K yield strength and fracture toughness of

Fe-8Cr-43Ni-2Mo welds

• Study of the effect of molybdenum on the cryogenic mechanical properties

of welds in more detail

• Continuation of mechanical property testing in support of superconducting

magnet design and construction

213
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS

Weld Alloy

(designation)

Welding

Process Tensile CharpyImpact Fracture Toughness

Fatigue Crack

Growth Rate Report Volume

(page)

Austenitic Stainless Steels

Fe-1 3Cr-20Ni-9Mn-2Nb SMAW I, II I, II, VIII I (159,169)

II ( 299) VIII (29)

Fe-15Cr-12Ni-10Mn SMAW IV IV, VIII IV (453) VIII (29)

Fe-1 6Cr-1 5Ni-HMo SMAW IV IV IV (275)

Fe-1 6Cr-l 6Ni-2Mo FCAW III III III, VI III (155)

VI (199)

Fe-1 6Cr-1 6Ni-ilMo FCAW IV IV, VIII IV IV (275) VIII (29)

Fe-1 6Cr-35Ni-ilMn SMAW I I, VIII II I (159)

(330 modified) II (299) VIII (29)

Fe-1 7Cr-1 6Ni-UMo SMAW IV IV IV (275)

Fe-l8Cr-l6Ni-9Mn-N GMAW III, IV IV III , IV III, VI III (155)

IV (275,453)

VI (199)

Fe-1 8Cr-l8Ni-2Mo-5Mn-N SMAW IV IV IV (453)

Fe-1 8Cr-20Ni-5Mn GMAW IV IV IV (275)

Fe-1 8Cr-20Ni-6Mn-0.3Nb GMAW IV IV IV IV (275)

Fe-1 8Cr-20Ni-6Mn-0.2Ti GMAW IV IV IV IV (275)

Fe-1 9Cr-1 2Ni-2Mo SMAW II II , VIII II II (315)

(316) VIII (29)

Fe-1 9Cr-1 2N1-2MO ESW IV IV IV, VI IV (275,415)

(316D VI (199)

GMAW III , IV III,IV,VIII IV III (155,195)

IV (275) VIII (29)

GTAW III, IV III, IV, VIII IV III (155,195)

IV (275) VIII (29)

SAW III, IV,

V

III ,IV,V III, IV III, VI III (155,195)

IV (275,415)

V (233)

VI (199)

SMAW II, III, IV II , VIII II, III, IV III , IV, VI II (299,315)

III (155,167)

VI (199) VIII (29)

* MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-88^1 (I); NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-16A5 (IV); NBSIR 82-1667 (V); NBSIR 83-1690 (VI); NBSIR 811-3000 (VII); NBSIR 85-3025 (VIII).

214
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS, continued

W- Id Alloy Welding Tensile Charpy Fracture Fatigue Crack Report Volume

designation) Process Impact Toughness Growth Rate (page)

Fe-19Cr-12Nl-2Mo-N SMAW II II, VIII II II (299)

(316LN) VIII (29)

Fe-19Cr-l6Ni-7Mn GMAW IV IV IV (453)

Fe-20Cr-10Ni SMAW I I, VIII I (169)

(308) VIII (29)

Fe-20Cr-10Ni FCAW III III III, VI III (155)

(308L) VI (199)

GMAW III III III, VI III (155)

VI (199)

SMAW II, III:, II, VIII II, III III, VI II (299)

III (155)

VI (199) VIII (29)

Fe-20Cr-10Ni-N FCAW IV IV IV, VI IV (275,415)

(308LN) VI (199)

Fe-20Cr-10Ni-9Mn-N SMAW I I I (159,169)

Fe-20Cr-l6Ni-7Mo-2W SAW IV IV IV (275)

Fe-20Cr-3'4Ni-2Mo-Nb SMAW VII VII VII (253)

Fe-21Cr-6Ni-9Mn-N SMAW I, II I, II I (159,169)

II (299)

Austenitic Steels

Fe-15Mn-8Ni-lMo-0.7C SAW VIII VIII VIII (15)

Fe-5Cr-25Mn-l Ni V V,VII V,VI V (233)

VI (199)

VII (245)

Fe-6.5Cr-25Mn-3Ni-TMo SMAW V V V,VI V (233. 24‘^. VI (199;

Aluminum Alloys

Al-5Mg GMAW III IV IV III (15= . 1 (5183) IV 3^ 5 ^

Al-5Mg-0.12 Ti GMAW III IV IV III ; n - , (5556) IV

* 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)

.

215

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Manganese and Nitrogen in

Stainless Steel SMA Welds for Cryogenic Service*'^

C. N. McGowan, T. A. Siewert, R. P. Reed

National Bureau of Standards, Boulder, Colorado

and

F. B. Lake

Alloy Rods Corporation, Hanover, Pennsylvania

Evaluation of a shielded metal arc weld test matrix in vhich

manganese (1.5 to 10 vit.%) and nitrogen (0.04 to 0.26 wt.^) were varied

independently has clarified the effect of these elements on the cryogenic

mechanical properties and predicted ferrite number (FN). Several molybdenum

and boron additions were also made, but they had no observable effect on strength or Charpy V-notch ( CVN ) absorbed energy. The matrix was based on a type 308L stainless steel weld metal composition. Desired compositions and

constant FN were attained through alloy additions to the electrode coating.

For each weld, one all-weld-metal 4-K tensile specimen and five ?6-K CVN

impact specimens were tested.

The addition of nitrogen linearly increased the 4-K yield strength from

600 to 1300 MPa and decreased the 76-K lateral expansion from 0.7 to 0.1 mm. Nitrogen reduced the 76-K CVN absorbed energy but not linearly. The addition

of manganese slightly increased the yield strength and slightly decreased the

lateral expansion. The 4-K tensile strength was relatively unaffected by

alloy additions; values varied between 1300 and 1500 MPa.

The DeLong FN predictive equation was improved by substituting a constant manganese term of 0.35 into the DeLong nickel equivalent. Analysis

of the remaining FN deviation revealed an interaction between nitrogen and

manganese. Addition of manganese-nitrogen interactive terms to the nickel

equivalent increased the accuracy of the FN prediction.

INTRODUCTION

Toughness values for 300-series stainless steel welds at cryogenic

temperatures have consistently been lower than those for base materials of

equal strength [1,2]. Several studies have attempted to understand this

phenomenon by statistically determining predictive equations for Charpy

V-notch (CVN) toughness on the basis of the weld composition [3“5]. Although

these studies have resulted in tougher welds through optimized electrode com- positions, the welds still do not match the toughness of the base material.

Recently, controlled additions of nitrogen (along with additions of manganese

* Work supported by the U.S. Department of Energy; not subject to copyright.

Submitted to Weld. J.

217
to reduce the tendency to form porosity) have increased the plate strength at

cryogenic temperatures with little loss in toughness [6]. However, this

increased base material strength cannot be totally utilized until welds of

comparable toughness at these elevated streng-ths are developed.

This study was intended to contribute to the development of welds with

properties comparable to those of the nitrogen-strengthened base materials.

We investigated the strength-toughness relationship in a series of shielded

metal arc (SMA) welds. The nitrogen and manganese contents were varied to

evaluate their effect on the weld properties. Several welds with varying

boron and molybdenum contents were also included.

FACTORS INFLUENCING STRENGTH, TOUGHNESS, AND FERRITE NUMBER

Strength

The strengthening potential of nitrogen in stainless steel increases

with a decrease in temperature [5,7-11]. Nitrogen occupies interstitial

sites and expands the lattice in a manner similar to carbon [12]. Unlike

high carbon contents, high nitrogen contents do not result in sensitization

problems. The nitrogen content can be increased until porosity occurs at the

solubility limit [9]. It can be introduced to the deposit through shieldinggas mixtures with the gas metal arc (GMA) or gas tungsten arc (GTA) processes

or through nitrogen-containing alloy additions to the coating with the SMA

process. With 0.05 wt./S nitrogen additions in type 308 stainless steel

welds, increases in yield strengths of approximately 120, 150, and 175 MPa

were reported by Mukai et al. (Ill K) , Enjo et al. (76 K) , and Onishi et al.

(76 K), respectively [13,10,11]. Onishi et al. report a similar Increase in

yield strength at 4 K. In comparison, only slight increases in weld yield

strength (< 50 MPa) have been found at room temperature [10,11,13,14].

The influence of nitrogen on the tensile strength is disputed. Onishi

et al. [11] and Ozeki and Ogawa [15] found increases in the 76-K tensile

strength of 100 and 400 MPa when nitrogen content was increased from 0.05 to

0.15 and 0.20 to 0.40 ut.%, respectively, whereas Enjo et al. reported a 76-K

tensile strength increase of only 20 MPa when the nitrogen content was raised

from 0.05 to 0.15 v^t.% [10]. In all cases, the magnitudes of tensile

strength increases reported were less than those observed for the yield

strength. The influences of other elements (such as manganese, molybdenum,

and boron) on cryogenic weld strength have not been published, but Sakamoto

et al. reported that manganese increased yield strength in stainless steel

base material at 76 K [16].

Toughness

The cryogenic toughness of type 308L and 31 6L stainless steel welds is

determined by the interrelated and sometimes opposing effects of the

individual alloying elements and the ferrite number (FN). The FN has been

found to be the dominant factor in determining CVN toughness in these

materials [4,17-20]. Lower FN tends to raise the toughness. On the other

hand, increased nitrogen contents, which lower the FN, have been reported

to lower the CVN absorbed energy and lateral expansion (LE) values in SMA

218
welds [^]. Contradicting this reported nitrogen effect, Enjo et al. [11]

found that CVN energies increased with increasing nitrogen content (decreasing FN) in his GMA study of type 308 stainless steel welds. These differences and the variations in toughness values near 0 FN indicate that the role

of nitrogen (and other alloying elements) with respect to toughness requires

further study. Several researchers [^,20] have studied the effect of

alloying elements that influence the ferrite content on the absorbed energy

at 76 K. Nickel reportedly increased toughness, whereas carbon and chromium

reduced toughness. Nitrogen is generally believed to reduce toughness, but

it has been suggested that the effect of nitrogen on CVN toughness may be

dependent on the specific alloy composition [19]. Manganese has been reported

to decrease 77~K CVN energy in stainless steel base material [16]. The

effects of molybdenum are not clear.

Ferrite Number

The strong influence of FN on cryogenic toughness dictates the use of an accurate method for predicting FN in the new high nitrogen and manganese

stainless steel electrodes. The diagrams developed by Schaeffler [21] and

DeLong [22] are currently the best methods available for the prediction of

FN . These diagrams were not designed for use with high manganese alloys;

they tend to underestimate FN for higher alloy contents because the aus- tenitizing power of manganese decreases at high manganese concentrations

[23,24]

.

Whether the various FN diagrams or the corresponding mathematical

expressions for these diagrams is used, nickel and chromium equivalents

represent the relative austenitizing and ferritizing tendency for a given

alloy composition. The DeLong nickel and chromium equivalents are:

Ni = Ni + 30 (C + N) + 0.5 Mn

;

Crgq = Cr + Mo + 1 .5 Si + 0.5 Nb

;

where the elements are in vit.%. Numerous studies have determined the

dependence of FN on alloying elements. The coefficients and constants

proposed to model the effects of nitrogen and manganese are given in table 1. At high nitrogen contents. Espy, in particular, found that the nitrogen coefficient decreases [24]. Most studies involving high manganese contents have

proposed constants to represent the effect of manganese and replace DeLong's

0.5 Mn term. The Hull study [25] was the only high manganese matrix to yield

a manganese coefficient. However, its second-order term does reduce the

influence of manganese on the nickel equivalent at high concentrations. A review of the alloy ranges used in the various studies did not disclose the

reason why different terms (either constants or coefficients) were proposed

to model the influence of the manganese and nitrogen in the nickel equivalent

express ion.

MATERIALS AND PROCEDURE

To clarify the role of manganese and nitrogen on the cryogenic

toughness, and FN , and to understand the reason for the varying and

conflicting effects reported in the literature, a test matrix based

streng th

,

some t im- on a typ'

219
308 stainless steel weld composition was evaluated. Alloy additions were

made to the SMA electrode coating. Manganese (1.5 to 10 and nitrogen

(0.04 to 0.26 wt./J) were varied independently within the matrix. Several

molybdenum (2 and 4 wt.5^) and boron (0.006 and 0.015 wt.^) additions were

also made to the matrix.

The weld matrix was designed with constant FN (near 3 FN) to remove

ferrite effects from the analysis. The FN of weld pads and the welds were measured magnetically. The measurement device was calibrated with ANSI/AWS

Standard 'A4.2-74.

The weld metal carbon and nitrogen were determined with individual

element analyzers; the phosphorus and sulfur, by an optical emission

spectrophotometer; and the remaining elements, by conventional analytical

techniques. The chemical composition and measured FN for the twenty-four

welds included in this study are listed in table 2. Molybdenum interference

in the optical emission spectrograph probably caused the variations in

measured phosphorus contents.

Test welds were produced using a series of 3.2-mm (1/8-in) diameter

electrodes that had core wires from a single heat of type 308 stainless

steel. The base material was 1 3“mm-thick by 305-mm-long mild steel plate.

To overcome the effect of dilution by the base material, the exposed faces

and backing strip were buttered with two layers of weld metal prior to

beginning the test weld, as specified in AWS A5.4-78. The same welder, power

supply, and welding parameters (110 A, 22 V , an identical bead sequence, and

a heat input of approximately 0.6 kJ/mm) were used to minimize weld variations. The interpass temperature was maintained at 93°C.

Two all-weld-metal, 6-mm-d iameter tensile specimens, oriented along the

axis of the weld, were machined from each weld. The second specimen,

reserved as a spare, was tested in a number of cases to check the repeatability of the results; in these cases, the average value is reported.

Tensile specimens were tested at 4 K (liquid helium) with a strain rate of

0.25 mm/min. A 25-mm gauge length was used. Specific details on the 4-K

tensile testing equipment and procedures have been reported previously [29].

Five 10-mm^ CVN specimens, with their notches oriented perpendicular to the

plate surface and located along the weld centerline, were removed from each

weld and tested at 76 K.

RESULTS AND DISCUSSION

Strength

The yield and tensile strengths determined for the various welds at

4 K are given in table 2 and plotted for the various nitrogen contents of the

welds in figure 1 . The yield strength increased linearly with the nitrogen

content. The equation relating the yield strength to the nitrogen content

was

Weld Yield Strength (MPa at 4 K) = 400 + 3700 (wt.^N). (1 )

220
The F value was 256 and the coefficient of determination, (a measure of the

quality of the regression), was 0.92. This equation corresponds closely with

a previously determined equation relating the yield strength to nitrogen

content for stainless steel plate [6]:

Plate Yield Strength (MPa at K) = 350 + 3^00 (vit.% N). (2)

The tensile strength remained near the average of 1400 MPa for nitrogen

contents up to p.l6 wt.% and averaged 1500 MPa for the three alloys that had

about 0.26 wt.)S nitrogen. This slight increase in the tensile strength was

not statistically related to the nitrogen contents.

Yield strength as a function of the weld manganese content is shown in

figure 2. The nitrogen contents were coded, since they are the dominant

streng thener . The 0.05, 0.10, and the 0.15 wt.^ nitrogen concentrations all

show approximately a 20-MPa increase in yield strength for every 1 wt.^

increase in manganese content. The equation including both the nitrogen and

manganese effects on yield strength is:

Weld Yield Strength (MPa at 4 K)

= 360 + 3400 (wt.5J N) + 14 (wt./o Mn) . (3)

The F value was 154 and the value was 0.94. There was some evidence that

the strengthening tendency of manganese changes as a function of nitrogen

content (see figure 2), but linear terms express the relative effects of

manganese and nitrogen contents quite well for the scope of this study.

Toughness

Figure 3 shows that the CVN absorbed energy at 76 K remains relatively

constant at 40 J for nitrogen contents of less than 0.15 wt./S. Nitrogen contents above 0.16 wt.5S did cause a decrease in absorbed energy. At 0.25 ^t.%

nitrogen, a reduction of approximately 50^ was found. All compositions

followed this trend closely with the exception of the 3.84 wt.% molybdenum

alloy, which fell far below the trend. No effect of manganese on the CVN

energy was determined.

Unlike the CVN absorbed energy, lateral expansion values decreased

linearly as a function of nitrogen content (figure 4). A decrease in lateral

expansion of approximately 86^ occurred when the nitrogen content was raised

from 0.03 to 0.26 wt.5J. The manganese content also affected lateral expansion. The lateral expansion of alloy 12 is 0.1 mm less than that of alloy 11 in table 3 (both = 19 wt.^ Cr, 0.15 vjt.% N, 1.5 FN); we attribute this to its

significantly greater manganese content. The alloys with molybdenum additions generally fit the observed trends; however, the lateral expansion

values of alloys with combined high nitrogen, manganese, and molybdenum contents (particularly alloys 13 and 6) were low.

The CVN absorbed energy and lateral expansion are linearly related

for the common type 308L and 316L stainless steel compositions [3,4]. The

nonlinear relationship observed in this study between CVN absorbed energy and

221
lateral expansion was therefore unexpected, and its cause was not determined.

We suspect that oxygen content, inclusion density, ferrite content, or other

factors are responsible for limiting the impact energy to 40 J. If this is

indeed the case, the effect of nitrogen would become apparent only at nitrogen

contents high enough to lower the CVN absorbed energy below the 40-J shelf.

The nonlinear relationship found in this study indicates that it may not

apply to other alloys. It also indicates a need to re-evaluate the use of

the 76-K CVN test to screen material for 4-K toughness testing.

The strength-lateral expansion relationship for these alloys is

illustrated in figure 5. This relationship is better than that for alloy

content-lateral expansion and indicates that the strengthening mechanisms

(interstitial for the nitrogen and substitutional for the manganese) are controlling the lateral expansion. Relating the lateral expansion to the

strength by linear regression techniques produced the equation:

LE (mm at 76 K) = 0.95 - 6 x 10“^ [Weld Yield Strength (MPa at 4 K) ] . (4)

This equation had an F value of 344 and an value of 0.94.

Ferrite Number

Table 4 gives the difference between the value of FN measured magnetically

on the weld pads (FNP) and that calculated using the DeLong predictive equation for the various welds. The AFN, or deviation between the measured and

predicted FN, increases as nitrogen and manganese concentrations increase.

Three approaches were used to examine and minimize these deviations.

A graphical method was used to express the influence of manganese and

nitrogen on the nickel equivalent. In figure 6, the difference in nickel

equivalents (ANig ) is plotted versus the manganese and nitrogen contents [ 23 ]. The ^ign is a measure of the difference between the DeLong

diagram nickel equivalent, found using the measured FN and calculated chromium equivalent, and a calculated partial nickel equivalent in which either

the manganese or nitrogen term is deleted. The difference between the two

equivalents measures the effect of the manganese or nitrogen. The slopes of

the best fit lines through the manganese -ANi or nitrogen-ANigg graphs

define the respective manganese and nitrogen coefficients. In figure 6k for

example, the slope indicates a nitrogen coefficient of I 6 (R^ = 0.70). A

more significance nitrogen coefficient (R^ = 0.91) of 21 is obtained,

however, when alloy 22 is excluded from the data. Alloy 22 is excluded owing

to its deviation from the trend established by nearly identical compositions.

Also, by replotting all the data as the DeLong diagram nickel equivalent

versus nitrogen, a nitrogen coefficient of 21 is obtained ( R^ = 0.93)* A

nitrogen coefficient of 21 agrees well with those reported by Hull, Espy, and

Mel’Kumor (see table 1).

The data (in figure 6B) indicate a slightly negative manganese coefficient. A negative slope has not been proposed previously; but this concept

is also supported by the Szumachowski-Kotecki type 309 Nig data [23]. It is

clear that the Espy, Szumachowski-Kotecki, and Hull terms and the line

222
representing this study all predict manganese effects better than the 0.5

DeLong manganese coefficient. If a constant manganese term is fitted, the

manganese constant (-0.25) is less than that proposed by either Espy (0.87)

or Szumachowski and Kotecki (0.35) [24,23].

Stepwise linear regression , the second method of analysis, was used to

determine the accuracy of existing predictive FN equations and to examine

possible improvements to these equations by changing the coefficients and

adding new variables. In this approach, twenty variations of the DeLong

[22,30], Schaeffler [21,30], and Seferian equations [31] were evaluated and

ranked in order of their significance during the stepwise regression. Within

each equation type, different combinations of nitrogen and manganese terms were used. For nitrogen, coefficients of 30, 24, 18, 13, and 0 were

included. For manganese, DeLong's 0.5 manganese coefficient and constant

manganese terms of 0.35 and 0.74 were included. One equation that contained

Hull's manganese terms was included. From this wide variety of predictive

equations and modifications, the best equation was selected for the manganese

and nitrogen ranges in this study.

In addition to these FN predictive equations, compositional terms such

p as Mn, N, Si, Mn“^, MnN, and ON were also available for selection by the

regression program. Selection of one of these terms following the selection

of a predictive equation indicated that its influence was not properly

modeled in the predictive equation.

Figure 7 shows the predicted values for three of these equations (FNCD,

FNCDK, FNCSK) plotted against the FN measured on the weld pad (FNP) and

illustrates the substantial differences between the predicted and measured FN

values. These equations are defined by:

FNCD = -30.65 + 3.49(Cr + Mo + 1.5 Si) - 2.5[Ni + 30(N + C) + 0.5 Mn]; (5a)

FNCDK = -30.65 + 3 49 ( Cr + Mo + 1.5 Si) - 2.5[Ni + 30(N + C) + 0.35]; (5b)

FNCSK = -39.1 + 43.5(Cr + Mo + 1.5 Si)/2.5[Ni + 30(N + C) + 2.35]; (5c)

where all elements are given in wt./^. As an example of the notation, FNCSK

indicates the Ferrite Number calculated using a Schaeffler diagram equation

modified to include nitrogen and using the Szumachowsk i-Koteck i manganese

constant (0.35). The D in FNCD represents the DeLong diagram FN equation.

The unmodified DeLong equation (FNCD) underestimated the measured FN by

as much as 24 FN when the nitrogen and manganese contents exceeded the ranges

for which the equation was proposed (figure 7A). Figure 7B shows the

improved FN predictions resulting from the replacement of DeLong's 0.5 manganese coefficient with the 0.35 constant manganese term (FNCDK). Although the

FN values generated from the FNCDK expression agree well with the measured FN

values at lower manganese and nitrogen content, overall the FNCSK equation

produces a more linear data set. When modified with a coefficient to adjust

its slope and given the appropriate intercept, FNCSK predicted the effect of

manganese and nitrogen most accurately. For this reason, FNCSK was always

223
chosen as the best predictor of FNP by stepwise linear regression analysis.

A similar result was found using the ferrite number measured on the actual

welds (FNW). Because the results using FNP or FNW are so similar, they will

be used interchangeably and are identified simply as FN . The equation found

through this method of analysis was;

FN = 3.0 + 0.26 (FNCSK). (6)

This equation had an F value of 98 and an value of 0.82. In each calculation, the program indicated the need for an additional nitrogen- or nitrogenmanganese-containing term in the predictive equation. Overall, this method

of analysis showed that although FN prediction is greatly improved by using

the modified Schaeffler equation, the manganese and nitrogen terms within the

FNCSK expression require further study.

Modifi cation of DeLong* s nickel equivalent was the third and final

approach to reduce the FN deviations. The nickel equivalent was calculated

from the chromium equivalent and measured FN . Only compositional variables

associated with the present DeLong nickel equivalent and boron were allowed

as choices during linear regression analysis with this method. Here, it was possible to determine which compositional terms were statistically significant and their respective coefficients. A revised nickel equivalent was found that better predicted FN in these high manganese-nitrogen alloys:

Nigq = Ni + 29(C + N) + 0.53(Mn) - 0.05(Mn)^ - 2.37(MnN) + 0.94(MnN)2 - 0.71; (7)

where all elements are represented in ^t.%. An F value of 220 and an value of 0.98 indicate a high statistical significance for the equation.

This modified nickel equivalent can be substituted for the DeLong nickel

equivalent to determine FN using the DeLong diagram, or it can be substituted

for the DeLong nickel equivalent in the mathematical representation of the

DeLong diagram, equation 5a, yielding:

FN = -30.65 + 3.49(Cr + Mo + 1.5 Si) - 2.5[Ni + 29(C + N) + 0.53(Mn) - 0.05(Mn)2 - 2.37(MnN) + 0.94(MnN)^ - 0.71]. (8)

In developing this equation, alloys with nitrogen contents greater than

0.22 wt.^ and FN of less than 1.0 were not Included. Since the DeLong

equation was shown to have a large scatter at very high nitrogen or manganese

contents in figure 6 and accurate measurements of FN near zero are difficult,

these limitations seem reasonable.

The manganese effect is decidedly dependent on the nitrogen level, as shown on the contour map of the equation in figure 8. At a low nitrogen

content, the nickel equivalent first increases as the manganese content is

increased. As manganese contents exceed the 4 to 5.5 wt.5& range, however,

the nickel equivalent begins to decrease. The observation of a change in the

austenitizing power for manganese within this concentration range has been

reported by Suutala [32], Hull [25], and Guiraldenq [33]. The fact that

increased nitrogen contents modify the behavior of manganese has not been

previously reported. At 0.20 wt.5& nitrogen for example, no characteristic

224
change in the nickel equivalent takes place as the manganese content is

increased. The nickel equivalent slowly and continuously decreases to lower

values. On the other hand, at 0.06 wt.^^ nitrogen, the nickel equivalent

increases from 12 to 12.7 and then decreases to 11.4 as the manganese content

is raised. A change in the nickel equivalent from 12 to 13 at the average

chromium content of 19.5 results in an FN change of 3 on the DeLong diagram.

If either of the MnN terms is removed from equation 7, the other becomes

insignificant, the magnitudes of their coefficients are reduced, and the

coefficient for nitrogen approaches 22. Such an equation is in agreement

with those found by Hull [25] and Mel'Kumor and Topilin [27]. It could be

argued that equation (6) is in fact an extension of those developed by Hull,

Mel'Kumor and Topilin and is also related to that found by DeLong [22] at

lower manganese levels.

In summary, the DeLong manganese coefficient of 0.5 is inaccurate for

these high manganese-nitrogen alloys. The replacement of DeLong's coefficient with the Szumachowsk i-Koteck i constant manganese term (0.35) improves

FN prediction, but it still does not completely describe the behavior of

manganese. The revised nickel equivalent proposed in equation 6, however,

improves FN predictions by quantifying the nitrogen-manganese interactions.

From the various manganese and nitrogen terms reported by researchers and the

seemingly contradictory nitrogen coefficients indicated through graphical

versus computer methods in this study, it is clear that interactions between

alloy elements must be more fully understood before better (and, one hopes,

simpler) FN predictive equations can be written.

CONCLUSIONS

1. The 4-K yield strength of these welds is primarily a function of the

nitrogen content. Manganese increases the strength slightly.

Both effects can be expressed by

Weld Yield Strength (MPa) = 360 + 3400 (wt.5^ N) + 14 {^t.% Mn). (3)

2. The 76-K lateral expansion had a more linear correlation to the strength

and the alloy content than the CVN absorbed energy. The lateral

expansion was expressed best by

LE (mm at 76 K) = 0.95 - 6 x 10~^ [Weld Yield Strength (MPa at 4 K)].

(4)

3. Present FN predictive equations are inaccurate for l8Cr-9Ni stainless

steels with high manganese and nitrogen levels. Two equations were

found to improve FN prediction for these alloys. The first is simpler;

the second is more accurate.

FN = 3.0 + 0.26[-39.1 + 43.5(Cr + Mo + 1.5 Si)/

2.5(Ni + 30(N + C) + 2.35)]. (6)

Ni = Ni + 29(C + N) + 0.53(Mn) - 0 . 05 (Mn)^ - 2.37(MnN) + 0.94(MnN)^

’ - 0 . 71 . ( 7 )

225
The second equation is a modified Nig^, which is to be used with the

DeLong diagram or equation (5a).

ACKNOWLEDGMENT

This work was supported by the Department of Energy, Office of Fusion

Energy.

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( 2 ): 413-430.

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226
steel at low temperatures down to K. Tetsu-to-Hagane 67 (13): 271.

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323- 334 . Cambridge, England: The Welding Institute.

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on stainless steel weld metal ferrite. Welding Journal 63 (5): 156s-l6ls.

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steel Welding Journal 6l (5): l49s-156s.

227
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.

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33. Guiraldenq, P. 1967. Science reviews of metals 64: 907-938.

228
Table 1 . Various Manganese and Nitrogen Coefficients Determined in Previous Studies

Researcher Coef f ic lent Prlnc iple Comments

or Constant Matrix Variations

N Mn (wt.%)

Schaeffler [21] - 0.5 Mn Ni (0 to 30) 308, 309, 310 , 316, 410, and 502

DeLong [22] 30 N 0.5 Mn N (0.03 to 0.22) 309, 308L, 316, 31 6L , 347 (Low Mn, N)

Hull [25] 18.14 N 0.11 Mn - Ni (0 to 22) 70 chill - cast stainless

0.0086 Mn2 Mn (0 to 20) steel alloys. Mo, Si, V,

N (0 to 0.15) W, Ti, Cb, Ta, AS,, Co, Cu,

C also varied. Cr = 14 to 20 wt.%

Espy [2^1] 30 (N-0.0145) 0.87 Mn (5 to 12.5) 4 Nitronic series alloys.

22 (N-0.045) N (0.13 to 0.33) N and Mn not varied systematically.

20 (N-0.145) The 3 N coeff. are for < 0.20,

< 0.25, and < 0.35 wt.% N, respectively

Hammar 114.2 N 0.31 Ni (9.0 to 114) 130 austenitic alloys. (Ingots)

and N (0.01 to 0.20) Mn content = 1 to 2 wt.%.

Svenson [26] C (0.014 to 0.10) Cr = 17 to 25 wt.^.

Mel'Kumor 20 N

- Ni (10 to 16) Austenitic Stainless Steel

and

Topilin [27]

N (0.05 to 0.46) 0.5 Mn, 25 Cr, 0.09 C

Szumachowski 30 N 0.35 Mn (1 to 12) 308 , 309L, 31 6L, 307

and

Kotecki [23]

Low nitrogen (0.05 wt.$)

Ok agawa 13.^ N 0.5 Mn N (0.04 to 0.29) GTA weld (no filler metal) on et al. [28] 304L base metal. N added through

shielding gas.

229
Table 2. Chemical Composition' and Ferrite Numbers of the Welds.

Alloy

No.

c Mn Si p s Cr Ni Mo N B FNPad FN Plate^

FN Calculated

DeLong

1 0,.033 1.57 0,,29 0 .013 0,.006 17.58 9.19 0. 02 0 .047 # 2.8 4.9 1.35

2 0,.03^ 1.1J9 0..34 0.,020 0,.006 15.37 9.06 2. 03 0..034 3.4 4.4 2.2

3 0-.03^1 3.23 0.,39 0,.014 0,.006 17 .87 9.02 0.,02 0 .035 # 4.0 4 .7 2.07

4 0.,036 6.5H 0.,37 0.,015 0..007 17.61 9.13 0. 02 0..046 2.8 4.2 - 4.3

5R 0,.039 6.31 0,,34 0

,

.021 0,.007 15.39 9.14 1

.

,99 0 .047 # 2.8 3.6 - 5.4

5 0,.039 6.31 0,,34 0,.021 0..007 15.39 9.14 1

.

99 0..047 # 2.8 3.6 - 5.4

6A 0,.036 6.27 0.,34 0..027 0..007 15.17 9.14 3.,84 0..162 # 2.6 2.9 - 8.06

7 0.,032 1.66 0.,32 0,,014 0,,006 18.36 8.94 0. 02 0.. 103 # 1.8 3.1 0.6

8 0,.033 2.96 0..36 0,.01 3 0..006 18.13 9.12 0. 02 0 .037 # 3.6 5.1 0.02

9 0..038 6.46 0..35 0,.015 0..007 18 .51 9.17 0. 02 0,,098 # 2.6 3.6 - 5.35

10 0..037 9.52 0,.38 0,.01 7 0,,009 17.71 9.10 0 .,02 0,.106 # 1.5 2.7 -12.16

1

1

0..031 3.18 0..34 0.,015 0..006 18.99 9.05 0. 02 0,. 153 # 1.4 1.9 - 2.93

12 0..037 6.62 0..38 0,.016 0,.008 19.51 9.06 0. 02 0..151 2.2 3.6 - 5.53

13 0,.035 6.61 0..32 0.,021 0.,007 17.77 9.10 1

.

66 0., 166 # 2.5 3.5 - 7.25

1A 0,.038 9.38 0..35 0,.017 0,.009 18.98 9.14 0. 02 0..163 # 1 .7 2.8 -1 2.16

15 0..032 6.11 0..36 0,,016 0.,007 20.90 9.29 0. 02 0,.252 # 1.6 2.0 - 7.92

16 0,.038 9.56 0..37 0,.019 0,.009 20.22 9.07 0. 02 0,.265 # 1 .2 1 .8 -1 4.75

17 0,.031 2.90 0. 30 0.,01 3 0.,006 18.79 9.12 0. 02 0.. 104 # 2.0 3.8 2.83

18 0,.033 3.25 0,.39 0,.014 0..006 17.58 9.00 0. 02 0,.032 0 .015 3.2 3.7 1.38

18R 0..032 3.16 0.,42 0,,01 3 0..006 17.90 9.05 0. 02 0..041 0,.006 3.0 4.6 2.04

19 0,.039 9.55 0,.34 0,.017 0,.009 17 .70 9.26 0. 02 0,.037 # 3.6 4.8 - 7.81

20 0..033 3.10 0..36 0..015 0,.006 17.75 9.07 0. 02 0,. 152 # 0.8 1.0 - 7.17

21 0,.033 3.46 0..37 0,.014 0,.006 17 .82 9. 10 0. 02 0,.150 0 .008 0.6 0.5 - 7.25

22 0,.032 9.72 0,.34 0.,018 0..009 17.60 9.06 0. 02 0., 259 0.2 0.7 -24.00

* Weight percent.

^FN plate is the average of FN bead surface and FN impact specimen surface.

^^Not determined.

230
Table 3. Mechanical Properties of the Welds

i(-K Tensile Properties 76-K CVN Properties

°y 0^ Elongation Reduction LE AE Alloy (MPa) (MPa) (?) (?) (mm) (J)

1 469 1586 26* 16.6 0.686 46

2 527 1570 20 13.5 0.71 1 39

3 504 1 432 25* 23.8 0.584 41

4

• 455 1341 26 22.7 0.584 42

5R 619 1346 27* 18.0 - -

5 665 1489 37 29.1 0.533 37

6 1169 1447 12 14.1 0.127 16

7 787 1412 21* 14.3 0.483 39

8 502 1387 24* 14 .9 0.686 43

9 864 1410 24* 22.4 0.432 42

10 871 1395 32 23.8 0.406 37

1

1

932 1376 20 16.8 0.406 40

12 1020 1376 18.8 15.3 0.305 37 13 1129 1410 12.0 13.5 0.229 33

1^1 1075 1527 28 .5 26.6 0.305 34

15 13^*1 1627 16.5 18.9 0. 127 20

16 1306 1622 20.0 13.3 0.127 24

17 748 1438 25* 10.2 0.508 42

18 502 856 6.5 10.8 0.610 43

18R 491 1267 19* 14.5 0.610 38

19 635 1299 27 23.1 0.483 37 20 888 1261 16* 15.1 0.457 42

21 874 1 1 25 11.3 9.6 0.432 41 22 1250 1498 15.5 15.4 0. 152 21

LE = lateral expansion; AE = absorbed energy.

* fractured outside the gauge marks

231
Table 4. Deviation between

the FN Calculated

the Measured

Using DeLong'

FN and

s Equation

Alloy wt . ? N wt.? Mn AFN*

1 0,047 1.57 1 .45

2 0.034 1.47 1.20

3 0,035 2.23 1 .93

4 0.046 6,54 7.10

5 R 0.047 6.31 8.20

5 0.047 6.31 8.20

6 0.162 6.27 10.66

7 0 . 103 1.66 1.20

8 0.037 2.96 3.58

9 0.098 6.46 7.95

10 0.106 9.52 13.66

1

1

0.153 3.18 4.33

12 0,151 6.62 7.73

13 0.166 6.61 9.75

14 0.163 9.38 13.86

15 0.252 6.11 9.52

16 0.265 9.56 15.95

17 0.104 2.90 - 0.83

18 0.032 3.25 1 ,82

18 R 0.041 3.16 0.96

19 0.037 9.55 11.45

20 0 , 152 3.10 7.97

21 0.150 3.46 7.85

22 0.259 9.72 24.20

*AFN is defined as the difference between the measured FN (FNP) and

the FN calculated with the DeLong predictive equation.

232
240

220

200

180

160

140

120

100

80

60

40

0 0.04 0.08 0.12 0.16 0.20 0.24 0.28

Nitrogen, wt. %

Fig. 1. Strength versus weld nitrogen content at 4 K.

1400

1300

1200 ^

1100 - 1000

^ 900

S 800

(U Ql

to 700

600

500

400

Q a o

D

•P

• 4 K Yield Strength

o 4 K Tensile Strength

233
76

K

CVN

Absorbed

Energy,

J

Strength,

MPa

200

180

160

140 -

120 ^

100

80

60

0 1 2 3 4 5 6 7 8 9 10

Manganese, wt. %

1400

1300

1200

1100

1000

900

800

700

600

500

400

1 1 1 1 1 1 —

—

V — V

V —

— V

V

V V • • — • • o

o

o —

- « o o 0.05 wt. % N

• O.lOwt. %N “

V 0.15 wt. % N

1 1 1 1 1 1 1 1 1 0.25 wt. % N

Fig. 2. Yield strength versus weld manganese content at 4 K.

50

45

40

35

30

25

20

15

10

5

0 0.04 0.08 0.12 0.16 0.20 0.24 0.28

Nitrogen, wt. %

Fig. 3. CVN absorbed energy versus nitrogen content at 76 K.

234

ft-lb
Fig. 4. Lateral expansion versus nitrogen content at /6 K.

in

0.010 0.020 0.030 in

76 K Lateral Expansion, mm

Fig, 5. Yield strength at 4 K versus lateral expansion at lb K.

235
Fig. 6. Difference in nickel equivalent versus weld nitrogen

content (A) and versus manganese content (B).

236
FNCSK

FNCDK

FNCD

FN

Fig. 7. Measured FN versus calculated FN.

237
Fig. 8. The nitrogen-manganese interaction expressed

in terms of the nickel equivalent.

238

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WS-;
STRENGTH-TOUGHNESS RELATIONSHIP FOR AUSTENITIC

STAINLESS STEEL WELDS AT 4 K*t

R. L. Tobler, T. A. Siewert, and H. I. McHenry

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

For the past decade the National Bureau of Standards has been compiling

cryogenic mechanical property data on structural alloy welds proposed for 4-K

service. Analysis of this data demonstrates that: 1) there is an inverse

correlation between yield strength and fracture toughness for stainless steel

welds at 4 K and that 2) the welds have significantly lower toughness than

base materials of comparable strength. The toughness of types 308L and 31 6L

welds is described by Kjq = 270 - 0.1 6 Oy.

INTRODUCTION

For the past decade the National Bureau of Standards has been compiling

cryogenic mechanical property data on structural alloy welds proposed for

magnetic fusion energy components designed to operate at liquid helium

temperature, 4 K. A sizable data base has been accumulated and certain

general trends are becoming apparent. The most distinct of these is the

correlation between yield strength and fracture toughness.

An adequate combination of yield strength (oy) and fracture toughness

(Kjq) makes a material suitable for structural applications. At any

temperature, however, these properties are antithetical. For wrought

austenitic stainless steel materials, Kjq estimates are generally found to be

a monotonically decreasing function of ay [1]. For a controlled series of

conventionally processed wrought AISI 304 type ( Fe-1 8Cr-1 ONi ) austenitic

stainless steels with varying carbon and nitrogen contents, the empirical

relationship is given as [2]:

Kjc = 500 - 0.3 Oy. (1

)

where is the fracture toughness in MPa*m^/2 j^g yield strength

in MPa. This equation has been used as a practical means of comparing and

evaluating current materials capabilities relative to projected goals or requirements for future superconducting magnet designs [3,4]. A corresponding

relationship for austenitic steel welds at 4 K, developed from the cryogenic

mechanical property data base, is presented and discussed in this paper.

^Contribution of the National Bureau of Standards, not subject to copyright.

tSubmitted to Cryogenics.

239
MATERIALS AND PROCEDURES

Thirty-three welds fabricated using various candidate electrodes and

welding processes have been tested at NBS at 4 K during the past ten years

[ 5

- 13 ]. Tensile tests were typically performed using cylindrical specimens,

6.35 mm in diameter and 25.4 mm in gage length. Plane strain fracture

toughness tests typically used compact or bend specimens 25.4 mm thick. In

most cases, J-integral tests were performed and Kjq was estimated from

measurements of Jjq [2]. A complete list of the welds is given in Table 1 along with values of Oy data and Kjq estimates and appropriate references.

The welds can be divided into two groups on the basis of composition.

The first group includes type 308 L and 31 6L weld compositions that are commonly used to join type 304 and 316 austenitic stainless steel base metals

for cryogenic structures. These welds have broad but clearly specified

compositional ranges so that their mechanical properties can be varied

significantly by compositional shifts within the allowable ranges. The second

group is unrestricted with respect to composition and includes some high

manganese welds. Both groups had a range of nitrogen contents so the strength

and toughness relationship could be evaluated over a range of strength.

The weld compositions determine the microstructures, which comprise two

groups also: single-phase (fully austenitic) or two-phase (austenite plus

delta ferrite), with ferrite number (FN) up to 12. The type 3O 8 L and 31 6L

weld compositions typically contain a small amount of ferrite to prevent

fissuring. The unrestricted compositions are more resistant to fissuring and

do not contain ferrite.

In addition to composition, the type of welding process used can significantly influence the purity and/or microstructure of the welds. The

various welding processes represented in this study are: shielded metal arc

(SMA), gas metal arc (GMA), gas tungsten arc (GTA), flux cored arc (FCA),

submerged arc (SA) and electroslag (ES).

RESULTS

Figure 1 is a plot of the fracture toughness versus tensile yield

strength at 4 K for the thirty-three welds of Table 1 . Despite wide

variations in composition and processing and sizable scatter in the data, an inverse correlation between yield strength and toughness is indicated. If one considers only type 3 O 8 L and 31 6L welds, the scatter is reduced and the

correlation becomes more apparent.

Figure 2 is the plot of yield strength versus fracture toughness

developed exclusively for type 308 L and 31 6L compositions. The data show no distinguishable effect of welding processes upon the strength-toughness

correlation, apart from the general scatter. Welds with FN > 7 show

relatively low toughness (these are identified on the figure) but welds with

lower FN contents fall randomly within the scatterband. Using a linear

regression technique, the following empirical relationship was found for the

welds of Figure 2:

240
Kjc = 270 - 0.16 Oy, ( 2 )

where Kjq is the fracture toughness in MPa*m''^2 ^ is the yield strength

in MPa. The one-standard-deviation scatterband of ±44 MPa*m^/^ is shown in

the figure.

In Figure 3 the one-standard-deviation scatterband for the welds (types

308L and 31 6L) and the trend line for the conventional base material are compared [2]. The Oy vs Kj^ relationships are dissimilar, with the welds

falling substantially below (about 40^) the base metals in performance.

DISCUSSION

Mechanical property data on 33 welds tested at 4 K indicate that the

stainless steel welds have a significantly lower toughness than similar base

metals of comparable strength. This is shown in Figure 3 where the one- standard-deviation scatterband for the welds (types 308L and 31 6L) and the

trend line for type 304 base metal [2] are compared. The differences in

toughness are particularly large at low yield strength.

Recently, five new wrought alloys have been developed in Japan [3,^1*

These stainless steel alloys are strengthened by nitrogen (0.2 to 0.35 wt.?),

like type 304LN, but are substantially different in chromium, nickel, and

manganese contents. The new steels include 25Cr-l 4Ni-4Mn, 1 2Ni-l 2Cr-l 0Mn-5Mo

,

1 3Cr-5Ni-22Mn , and 1 5Cr-l Ni-25Mn. With yield strengths and fracture

toughnesses of 1200, MPa and 200 MPa*m’''^2^ respectively, these alloys are superior to the conventional base metals in Figure 3. Thus, the gap between

base metal and weld metal performance continues to wicjen. This gap emphasizes

the practical need for further development of filler metals.

The strength- toughness data obtained from stainless steel welds are from

a wide range of compositions, processes, and welding procedures. Yet none of

the data are near the Japanese target values of 1200-MPa yield strength and

200-MPa*m‘/^ fracture toughness. Below, we consider the causes of the

relatively poor toughness of the welds.

Since welds are metallur gically more complex than wrought base metals,

the upgrading of weld properties represents a significant challenge. To

achieve properties comparable to those of the base metals, the strengthtoughness characteristics of welds at 4 K must be improved by a substantial

margin.

Austenitic stainless steel welds generally fail by a ductile fracture

mechanism, i.e., by the formation and growth of voids that eventually compose

the fracture surface. Thus, toughness improvements must increase the

resistance of the weld metal to the nucleation and growth of voids. Voids

nucleate most readily near phase boundaries, as a result of interfacial

separation, fracture within the second phase, or matrix separation caused by

strain concentration. Toughness improvements would be expected in stainless

steel welds if the other phases, such as delta ferrite, chromium carbides, and

inclusions, were eliminated or reduced. Previous studies have demonstrated

the deleterious effects of a multiphase structure.

241
A study of four type 31 6L welds with an FN range of 0 to 12 has shown a clear trend toward lower toughness with increasing FN [7]. With the large

scatter caused by differing weld processes and techniques in this study, the

effect of FN is less apparent. But when the FN is greater than 7, as labeled

in Figure 2, the data clearly fall along the lower edge of the scatterband.

Thus, the ferrite should be reduced to the lowest level consistent with

fissure resistance.

The deleterious effect of sensitization on weld metal toughness [7] is

due to the precipitation of chromium carbides, which serve as void initiation

sites. The influence of inclusions on weld toughness has been studied

indirectly by Whipple and Kotecki [12] who produced a series of type 31 6L

welds using three different processes: GTA, GMA, and SA. The toughness at

4 K was inversely related to the inclusion content, with the highest toughness

in the GTA welds (I8l MPa*m^/^) and the lowest toughness in the SA welds (104

MPa *m .

The growth and coalescence of voids relates to the strength, ductility,

and strain-hardening behavior of the matrix. As matrix strength increases,

less energy is dissipated by plastic deformation during void growth and,

consequently, toughness is reduced. Increased matrix strength also tends to

activate additional void nucleation sites. Thus, increasing matrix strength

by alloying with interstitial elements such as carbon and nitrogen tends to

reduce fracture toughness, and the inverse relationship between strength and

toughness shown in Figure 3 results. One exception to this inverse relationship may be the influence of grain size as shown by Whipple, McHenry, and Read

[9]. As the width of the columnar grains decreased in type 31 6L SMA welds,

the yield strength changed slightly (800 MPa ± 1055), but toughness increased

significantly (from 95 to 200 MPa*m^^^).

In summary, to increase the strength-toughness characteristics of welds,

the resistance to ductile fracture must be increased. Analysis of existing

data indicates that this can be done by eliminating delta ferrite, reducing

the inclusion content, avoiding precipitates such as chromium carbides, and

reducing the width of the columnar grains. These approaches are not expected

to eliminate the inverse relationship between strength and toughness, but to

raise the trend line to a level approaching that achieved in the wrought

stainless steels.

GONGLUSIONS

This study demonstrates that 1 ) there is an inverse linear correlation

between yield strength and fracture toughness for stainless steel welds at

4 K, and that 2) the welds have significantly lower toughness than base

materials of comparable strength.

ACKNOWLEDGMENT

This research was sponsored by the Office of Fusion Energy, with

Dr. V. Der as project manager.

242
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1. Read, D. T. , Reed, R. P. Cryogenics, 2j_ (1981) 415. [see also;

Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures - II, NBSIR 79-1609, National Bureau of Standards, Boulder,

Colorado (1 979) 79-122].

2. Tobler, R. L. , Read, D. T. , Reed, R. P. Fracture Mechanics: Thirteenth

Conference, ASTM STP 743, R. Roberts, ed

.

, American Society for Testing

and Materials, Philadelphia, (1981) 250.

3. McHenry, H. I. Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures - VIII, NBSIR 85-3025, National Bureau of Standards,

Boulder, Colorado (1 985 ) 323.

4. Shimamoto, S. , Nakajima, H. , Yoshida, K., Tada, E. Advances in Cryogenic

Engineering—Materials, vol. 32, Plenum (1986) 23.

5. McHenry, H. I. Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures - I, NBSIR 78-884, National Bureau of Standards,

Boulder, Colorado (1 978) 159.

6. McHenry, H. I., Read, D. T. , Steinmeyer, P. A. Materials Studies for

Magnetic Fusion Energy Applications at Low Temperatures - II, NBSIR

79 - 1609 , National Bureau of Standards, Boulder, Colorado (1979) 299.

7 . Read, D. T. , McHenry, H. I., Steinmeyer, P. A., Thomas, R. D., Jr.

Welding Journal ^ (1980) 104s.

8. McHenry, H. I., Whipple, T. A. Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures - III, NBSIR 80-1627, National

Bureau of Standards, Boulder, Colorado (1980) 155.

9. Whipple, T. A., McHenry, H. I., Read, D. T. Welding Journal ^ (1981)

70s.

10. Whipple, T. A., McHenry, H. I. Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures - IV, NBSIR 81-1645, National

Bureau of Standards, Boulder, Colorado (1981) 273.

11. Elmer, J. W., McHenry, H. I., Whipple, T. A. Materials Studies for

Magnetic Fusion Energy Applications at Low Temperatures - IV, NBSIR

81-1645, National Bureau of Standards, Boulder, Colorado (1981) 289.

12. Whipple, T. A., Kotecki , D. J. Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures - IV, NBSIR 81-1645, National

Bureau of Standards, Boulder, Colorado (1981) 303.

13 . McHenry, H. I., Elmer, J. W., Inoue, T. Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures - V, NBSIR 82-1667,

Natinal Bureau of Standards, Boulder, Colorado (1982) 245.

243
Table 1 . Weld Compositions and M-K Mechanical Property Data*

Designation Welding Composition (weight percent) FN 4-K Properties Ref. Process Oy *^Ic 1/2 (MPa*m ' C Mn Si P S Cr Ni Mo N (MPa) :

1

330 M SMA 0.14 4.6 0.48 0.019 0.006 17.8 34.4 0.3 0.04 0 911 153 5 308L-15 SMA 0.02 1.9 0.22 — — 19.5 10.6 0.1 0.06 4.1 654 156 6 316L-15 SMA 0.03 2.0 0.30 — — 17.5 13.7 2.1 0.04 0.2 769 118 6 316L-15 SMA 0.03 1.9 0.32 — — 17.5 13.6 2.1 0.07 0 879 95 6 316L SMA 0.03 1.7 0.58 0.033 0.011 18.5 14.0 2.3 0.04 0.5 694 179 7 316L SMA 0.03 1.7 0.56 0.034 0.012 18.6 13.0 2.3 0.05 3.5 798 141 7 316L SMA 0.03 1.7 0.57 0.034 0.009 19.2 12.0 2.3 0.04 9.0 886 108 7 316L SMA 0.03 1.8 0.59 0.029 0.008 19.7 11.0 2.2 0.07 12.0 881 98 7 316 SMA 0.06 1.8 0.56 0.033 0.009 18.7 12.9 2.3 0.04 0.6 924 154 7 308L FCA 3.0 1117 66 8 308L GMA 7.0 776 167 8 316L SA 4.7 918 132 8 316L SA 8.0 860 99 8 316L SMA 882 171 9 316L SMA 727 199 9 316L SMA 861 132 9 316L SMA 815 161 9 316L SMA 772 200 9 316L SMA 845 187 9

31 6L ES 690 189 10 308L FCA 558 134 10 20-16-7 Mow SA 0.03 6.7 0.09 0.010 0.003 19.5 16.1 2.8 0.23 0 1110 148 11 17-16-J4 Mo SMA 0.03 4.0 0.39 0.022 0.007 17.3 15.9 2.2 0.05 0 814 250 11 18-20-6 Nb GMA 0.04 5.8 0.32 0.005 0.007 18.0 20.3 0.1 0.04 0 696 227 11 18-20-6 Ti GMA 0.03 5.7 0.28 0.006 0.007 18.2 20.3 0.1 0.02 0 686 287 11 18-16-9 GMA 0.05 8.6 0.44 0.004 0.021 18.1 15.8 0.08 0 1076 121 11 16-16-4 Mo FCA 0.03 4.4 0.25 0.012 16.1 15.9 1 .5 0.05 0 606 239 11 16-15-4 Mo GMA 0.02 3.7 0.22 0.020 0.012 15.7 15.0 1.8 0 717 170 11

31 6L GTA 0.02 1.4 0.45 0.020 0.003 18.0 11.9 2.1 0.07 1.5 807 181 12

31 6L GMA 0.02 1.4 0.47 0.020 0.003 18.0 11.8 2.1 0.07 3.0 779 163 12 316L GMA 0.01 2.2 0.46 0.010 0.01 3 18.7 15.2 2.3 0.03 0.9 608 159 12 25-5-1 SA 0.19 25.0 0.63 0.010 0.003 6.5 2.8 1 .4 654 164 13 25-5-1 SMA 0.16 26.6 0.13 0.003 0.006 6.4 3.2 1 .2 962 158 13

Since these welds were obtained from various sources, the material characterization is not uniform.

244
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Figure 1 . Fracture toughness versus tensile yield strength at 4 K for weld

metals representing various processes and compositions.

Figure 2. Fracture toughness versus tensile yield strength at ^ K for type

308 and 316 weld metals in Figure 1. A one-standard-deviation

scatterband of ±M4 MPa*m^/^ is included.

245
Figure 3. Comparison of the weld and base metal strength- toughness

relationships

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Tensile and Fracture Properties of an Fe-1 8Cr-20Ni-5Mn-0 . 1 6N

Fully Austenitic Weld Metal at 4 K*"*"

C. N. McGowan, T. A. Siewert, and R. L. Tobler

National Bureau of Standards

Fracture and Deformation Division

Boulder, Colorado

The 4-K tensile and fracture toughness properties of a fully austenitic

stainless steel weld are reported. One tensile and two compact tension

fracture specimens were tested. The weld was produced by gas metal arc welding using an Fe-1 8Cr-20Ni-5Mn-0. 1 6N electrode and a 98 % argon-2^ oxygen

shielding gas mixture. The yield strength of 1015 MPa and average fracture

toughness, Kj^(j), of 203 MPa'/m are higher than those of welds produced with

308L and 31 6L electrodes and compare favorably with base metal properties.

Examination of the fracture surfaces of all samples by scanning electron

microscopy showed ductile failure by microvoid coalescence. The suitability

of this alloy for welding cryogenic structures is discussed.

INTRODUCTION

As part of the continuing search to find improved welding electrode com- positions for cryogenic applications, the 4-K tensile and fracture, toughness

of a high nickel-manganese-nitrogen stainless steel weld were measured; the

results are reported here and compared to the U.S. strength-toughness goals

for superconducting magnet structural alloy welds. These weld goals (1000

MPa yield strength and 150 MPa/m fracture toughness at 4 K) were established

to guide improvements in welding materials and procedures [1].

The weld metal solidification structure is single-phase, face-centeredcubic austenite. Types 308L and 31 6L austenitic stainless steel weld metal,

currently used in the fabrication of cryogenic structures, typically have 0.5

to 9 % delta ferrite present in their respective structures. Even though the

ferrite phase is brittle and exhibits reduced toughness at low temperature

[2], it has traditionally been tolerated in the microstructure of these

alloys because it reduces their hot-cracking susceptibility [3~5]. Recently,

however, some success has been reported in attaining fully austenitic

microfissure-free welds [6-8]. The fully austenitic Fe-1 8Cr-20Ni -5Mn-0 . 1 6N

Contribution of the National Bureau of Standards, not subject to copyright.

Submitted to J. Eng. Mater. Technol.

247
electrode composition evaluated in this study achieves its fissure resistance

through control of segregation by reduced phosphorus and sulfur concentrations. Limiting the concentrations of these elements reduces the grain

boundary liquation caused by the subsequent weld beads [9].

MATERIAL

A nine-pass gas metal arc weld was made using an Fe-1 8Cr-20Ni-5Mn-0 . 1 6N

stainless steel electrode. The chemical composition of the 1 . 2-mm-d iameter

electrode is given in Table 1 . The base material, type 30^ stainless steel,

had a 25-mm thickness and was beveled to provide a 60° included angle for the

weld joint. Figure 1 shows the weld joint configuration and specimen locations. The specimens were located on the weld centerline to allow evaluation

of the undiluted weld metal. The welding was done at 230 A and 22 V. The

travel speed, weave width, and electrode extension values were maintained at

250 mm, 7 mm, and 22 mm, respectively. A shielding gas composition of

98/S argon-2/S oxygen was used.

The macrograph in Figure 2 illustrates the dendritic crystal structure

present within the test specimens. Orientation differences between dendritic

crystals are highlighted by preferential attack of the acid etch. The dark

circular areas are welding flaws, which were present only at the start of the

weld. The dendritic crystals grow from the weld joint edges toward the weld

bead center and into the plane of the page following the welding direction.

The fracture specimen notches were positioned to encourage crack growth in

the direction of crystal growth (the direction of welding in Figure 1) to

provide the most conservative measurement of toughness. The fully austenitic

microstructure and the precipitate distribution of the weld are shown in

Figure 3. la this high-magnification micrograph, the long primary dendrites

and the short secondary arms characteristic of primary austenite solidification appear as lighter etched areas. The interdendni tic areas are etched

darker due to enrichment in chromium and depletion in nickel during the

solidification process.

EXPERIMENTAL PROCEDURES

The tensile and compact tension specimens were tested at 4 K following

previously developed procedures [10]. The AWS A. 5. 4 all-weld-metal tensile

specimen (6.3”mm diameter, 25~mm gauge length), was tested at a strain rate

of 4 X 10 ^ s The two compact tension specimens (25“mm thick, 50.8-mm

wide) were modified for loadline displacement measurements [8] and the notch

was oriented normal to the weld axis, as discussed previously. The compact

tension specimens were precracked at 76 K, then tested at 4 K. The maximum

stress intensity factor during fatigue at 76 K was 33 MPa*/nT, and the stress

ratio during fatigue was 0.1. The single-specimen unloading compliance

method was used to obtain fracture toughness values, and estimates of

Kic(J) for the test specimens [10].

After mechanical testing, the fracture surfaces were examined with a torsion bar magnetometer and a scanning electron microscope (SEM).

248
RESULTS AND DISCUSSION

A 1015-MPa yield strength was determined at 4 K for the weld material

(Table 2). As expected of welds having 0.16 wt.^^ nitrogen, this yield

strength meets the U.S. goal of 1000 MPa for superconducting magnet structural alloy welds [1]. The tensile specimen had a total elongation of 26^

and a reduction in area of 375^.

The Kj^(j) fracture toughness values of the two compact tension specimens were 206 and 200 MPa-/rn (Table 3)- The J-resistance curve of each

specimen consisted of four J-Aa points in the exclusion rule interval [11].

These values were determined from by the expression:

Klo(J) .

where E is assumed to be 207 GPa at 4 K. The Kj^(j) values for both specimens exceed the 150-MPa*/m fracture toughness goal for cryogenic structural

alloy welds with a 1000-MPa yield strength.

When the compact tension fatigue cracks were inspected, the crack fronts

were found to be somewhat irregular and bordered on unacceptability, according

to the criteria established by ASTM E 8l3"8l. This type of fatigue crack

front irregularity is not unusual for welds owing to residual stresses [12].

The first specimen had a cusp-shaped fatigue crack front but was acceptable

according to the crack curvature requirements. On the second specimen, one crack edge exceeded the other edge by 11^, whereas the criterion is 10/5 or less.

The fracture surfaces of the tensile and the two compact tension specimens were examined by SEM. Ductile dimpling was found to be the dominant

feature of the tensile specimen fracture surface (Figure 4). Upon examination of the compact tension specimens, ductile dimpling again dominated the

fracture surfaces, indicating that microvoid coalescence was the microscopic

failure mode for both types of specimens. The fractograph, shown in Figure 5,

contains areas representing the commonly observed surface features of the two

compact tension fractures: (1) equiaxed dimples which nucleated as a result

of plastic flow about precipitates, (2) small areas of shear dimpling on regions fracturing out of the primary fracture plane, and (3) secondary

cracking (shown at the top of the fractograph). All three features indicate

that the weld behaved in a ductile manner during testing at 4 K. No evidence

of microf issur ing or martensite was found through SEM examination or magnetometer inspection.

Although analysis of the compact tension data and fracture surfaces has

shown this weld composition to have very good toughness at 4 K, for an engineering material, it is the relationship between strength and toughness that

ultimately determines its usefulness. In Figure 6, fracture toughness versus

yield strength at 4 K is plotted. The trend line representing the type 304

stainless steel base material family is based on data from seven heats [10]

and includes alloys with a range of nitrogen contents. The band representing

weld properties is based on data from twenty-three type 308L and 31 6L stainless steel welds made with several different welding processes [13]. The

249
type 308L and 31 6L welds have toughness inferior to the type 304 base

material at all strength levels.

The strength-toughness datum for the Fe-1 8Cr-20Ni -5Mn-0 . 1 6N composition

is also shown in Figure 6, along with data for three other recently developed

fully austenitic compositions [7,8]. These weld compositions exhibit toughness superior to types 308L and 31 6L and are similar to those of the type 304

base material. The improved toughness of these new fully austenitic weld

compositions is attributed to a increase in the energy required for microvoid

nucleation and growth during ductile fracture.

Microvoid nucleation in a stainless steel with a single phase matrix

commonly occurs at inclusions. A polished section through the weld

(Figure 7) reveals an average of 1000 inclusions per 0.1 mm^. The fracture

surface of the Fe-1 8Cr-20Ni-5Mn-0 . 1 6N weld in Figure 4 contains an average of

12,000 void initiation sites per 0.1 mm^. The order-of-magnitude difference

between the measured inclusion density and void initiation site density

occurs because a planar surface is being compared to a three-dimensional

fracture surface.

CONCLUSIONS

The 1015-MPa yield strength and the 203-MPa*/m values determined for

the fully austenitic stainless steel weld at 4 K meet current weld goals of

the U.S. fusion energy program. It is, therefore, a good candidate for

cryogenic applications where high strength and toughness are required.

ACKNOWLEDGMENT

The welding electrode was provided by the Wire Division of Sandvik

Steel. This work is supported by the U.S. Department of Energy, Office of

Fusion Energy.

REFERENCES

1. McHenry, H. I., Strength and Toughness Goals for Cryogenic Steels, in:

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures - VIII, NBSIR 85- 3025 , National Bureau of Standards,

Boulder, Colorado, 1981, p 121.

2. Brown, E. L., Whipple, T. A., and Tobler, R. L. Fracture Toughness of CF8

Castings at Four Kelvin, Metall. Trans. A, Vol. 1 4A , June 1983, pp. 1179“

1187.

3. Arata, Y. , Matsuda, F., and Katayama, S., Solidification Crack Susceptibility in Welds of Fully Austenitic Stainless Steels (Report II) - Effect

of Ferrite, P, S, C, Si, and Mn on Ductility Properties of Solidification

Brittleness, Trans. JWRI , Vol. 6, No. 1, Welding Research Institute of

Osaka University, Osaka, Japan, 1977.

250
Matsuda, F., Nakagawa, H., Katayama, S., and Arata, Y. , Solidification

Crack Susceptibility in Welds of Fully Austenitic Stainless Steels

(Report VID-Effect of Mn and N on Solidification Crack Resistance, Trans.

JWRI , Vol. II, No. 2, Welding Research Institute of Osaka University,

Osaka, Japan, 1982.

5. Brooks, J. A., Microstructure and Cracking of Austenitic Stainless Steel

Welds, Dissertation, Dept, of Metallurgy and Material Science,

Carnegie-Mellon University, Pittsburgh, 1981 .

6. Haddick, G. T., Thompson, L. D., Parker, E. R., Zackay, V. F., New

Nickel-Free Austenitic Stainless for Ambient and Cryogenic Service, Met.

Prog., Nov. 1978, pp. 37-^0.

7. Elmer, J. W. , McHenry, H. 1., and Whipple, T. A., Strength and Toughness

of Fully Austenitic Stainless Steel Filler Metals at 4 K, in: Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures - IV,

NBSIR 81-1645, National Bureau of Standards, Boulder, Colorado 1981,

pp. 289-302.

8. Tobler, R. L., Trevisan, R. E., and Reed, R. P., Tensile and Fracture

Properties of an Fe-1 4Mn-8Ni -1 Mo-0 . 7 Fully Austenitic Weld Metal at 4 K,

Cryogenics, Vol. 25, 1985, pp. 447-451.

9. Lundin, C. D. and Spond, D. F., The Nature and Morphology of Fissures in

Austenitic Stainless Steel Weld Metals, Weld. J., Vol. 55, Nov. 1976,

pp. 356s-367s.

10. Tobler, R. L., Read, D. T., and Reed, R. P., Strength/Toughness

Relationship for Interstitially Strengthened AISI 304 Stainless Steels

at 4 K Temperature, in: Fracture Mechanics: Thirteenth Conference,

Richard Roberts, Ed., ASTM STP 743, American Society for Testing and

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

11. Standard Test Method for Jj^, A Measure of Fracture Toughness, E8l3~8l,

in: 1982 Annual Book of ASTM Standards, Part 10, American Society for

Testing and Materials, Philadelphia, 1982, pp. 822-840.

12. Towers, 0. L. and Dawes, M. G., Welding Institute Research on the Fatigue

Precracking at Fracture Toughness Specimens, in: Elastic-Plastic Fracture

Test Methods: The User's Experience, E. T. Wessel and F. J. Loss, Eds.,

ASTM STP 856, American Society for Testing and Materials, Philadelphia,

1985, pp. 23-46.

13- Tobler, R. L., Siewert, T. A., and McHenry, H. I., Strength-Toughness

Relationship for Austenitic Structural Steel Welds at 4 K, submitted to

Cryogenics.

251
Table 1 . Electrode composition, wt.^.

c Mn Si P S Cr Ni Mo N Cu Co AS, B

0.033 5.4 0.26 0.006 0.007 18.1 20.4 0.1 0.16 0.015 0.02 0.002 0.001

Table 2. Tensile test results at 4 K.

Specimen Yield Strength,

MPa

Tensile Strength,

MPa

Elongation,

%

Reduction

of Area, %

1 1015 1300 26 37

Table 3. Fracture toughness test results at 4 K.

Spec imen Thickness,

mm Width

,

mm

Crack Length,

mm Fracture

kJ/m^

Toughness

Kjc(J).

MPa/m

A 23.9 50.8 27.8 206 206

B 23.9 50.8 28.7 194 200

252
Fig. L. Weld joint design and specimen orientation.

Fig, 2, Macrograph of a transverse section of a weld made

with a Fe-18Cr-20Ni-5Mn-0. 16N electrode.

253
Fig. 3, Micrograph illustrating the austenitic microstructure

and precipitate distribution.

Fig, 4. SEM fractograph of the tensile fracture surface.

254
Fig. 5. SEM fractograph of compact tension fracture surface.

Fig. 6. Fracture toughness versus yield strength relationships

for a number of type 308L and 316L stainless steel welds,

austenitic steel welds, and type 304 base material.

255
Fig. 7. SEM micrograph of a polished section of a weld

made with the Fe-18Cr-20Ni-5Mn-0. 16N electrode

illustrating the inclusion size and distribution.

256
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TENSILE, FATIGUE, AND FRACTURE PROPERTIES OF AN

Fe-l8Cr-l6Ni-6.5Mn-2.4Mo FULLY

AUSTENITIC SMA WELD AT 4 Kt Trevisan,* R. L. Tobler, R. E. * H. I. McHenry, P. T. Purtscher,

and T. A. Si ewer

t

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

A fully austenitic stainless steel butt weld was produced by

shielded metal arc (SMA) welding using a new filler composition:

Fe-l8Cr-l6Ni-6.5Mn-2.4Mo. The weld was 25-mm thick and contained no

mi crof issures . Tensile, fatigue crack growth rate, and fracture

toughness properties for the weld at 4 K are reported. Fatigue

crack growth resistance compares favorably with base metals for the case of crack growth in the longitudinal direction of the weld. A

completely ductile fracture with very fine dimple size was evidenced

in both tensile and fracture toughness tests. The 4-K weld fracture

toughness (Kjq = 190 MPa*m^2) is about 20^ lower than typical

values for conventional austenitic stainless steel base metals of

equivalent yield strength ( Oy

= 842 MPa) but is significantly higher

than values for conventional type 308L and 31 6L SMA welds, which

contain some brittle fracture features owing to second-phase

6-ferrite. The lower toughness of the weld relative to base metals

is attributed to the effects of a higher density of inclusions on ductile fracture initiation.

Key words: austenitic stainless steel; cryogenic properties of

materials; ductile fracture; fatigue crack growth; fracture

toughness; microstructure; shielded metal arc welding; tensile

properties; weldments.

INTRODUCTION

A current goal of cryogenic research is to develop microfissure-free,

fully austenitic welds that match base metal levels of strength and toughness

at liquid helium temperature (4 K). Austenitic welds traditionally contain

some second-phase 6-ferrite to help prevent microf issur ing during cooling to

room temperature. Unfortunately, the 6-ferrite is undesirable for cryogenic

service; having a body-centered cubic structure, it is brittle at low

temperatures

.

tWork supported by the U.S. Department of Energy; not subject to copyright.

*Guest worker, on leave from the Department of Engineering, UNICAMP,

Campinas - Sao Paulo, Brazil.

257
To prevent formation of 6-ferrite and achieve a fully austenitic weld, an appropriate chemical composition must be selected according to the DeLong

constitution diagram. Other factors must then be considered, including

microf issures , fusion zone width, structural inhomogeneity, austenite-tomartensite transformation tendencies, and inclusion contents [1]. Additional

research is necessary to identify the optimum alloy composition and weld

process variables, but it does appear possible to obtain both microfissurefree and 6-ferrite-free welds having attractive properties at 4 K [2,3]. In

this paper we consider a shielded metal arc (SMA) Fe-l8Cr-l6Ni-6.5Mn-2.4Mo

weld composition; evaluate its tensile, fatigue, and fracture properties at

4 K; and compare its performance with that of other SMA welds [4] and

austenitic stainless steel base metals.

MATERIAL

The test material was an SMA butt weld produced by joining two sections

of 25 -mm-thick Fe-l8Cr-8Ni base metal with the newly developed Fe-1 8Cr-1 6Ni6.5Mn-2.4Mo electrode having a rutile coating. The chemical compositions and

welding conditions are shown in Tables 1 and 2. The microstructure was characterized by various techniques, including li^t microscopy, energy

dispersive microchemical analysis, and magnetometr y.

EXPERIMENTAL PROCEDURES

Tests at 4 K were performed in the as-welded condition using specimen

geometries with the nominal dimensions shown in Figure 1. The tensile

properties were measured using three cylindrical (6.35-mm diameter by 25.4-mm

gage length) specimens at a strain rate of 4.2 x 10"“* s" ^

. Fatigue crack

growth rates and J-integral fracture toughness values were measured using

compact specimens with compliance measurements at the specimen edge and

loadline, respectively. The fracture specimens were 25“mm thick with a width-to-thickness ratio (W/B) of 2. The fatigue crack growth rate specimens

were 6.4-mm thick with a W/B of 12. Specimen orientations relative to the

weld metal are shown in Figure 2. The cryogenic apparatus and procedures are described elsewhere [5“8].

Single-specimen J-integral tests were used to measure J 10 , which is a material parameter for ductile fracture [6]. The computer-aided

unloading-compliance technique [6-8] was applied to generate resistance

curves, J versus crack extension (Aa) measurements, for three specimens. Jjq was defined as the intersection of the resistance curve regression and

blunting lines (see the Results section), and estimates of the plane strain

linear-elastic stress intensity factor, Kjq(J), were calculated from the

relation [6]

( 1 )

The value of Young’s modulus, E, was estimated to be 207 GPa at 4 K.

258
The first of the three J-integral test specimens was fractured completely

at ^4 K, and then its fracture surface was viewed by scanning electron

microscopy (SEM). The second was tested and then metallographically sectioned

in the plane perpendicular to the crack surface to observe the crack profile.

The third was tested and then refatigued to measure the final physical crack

extension increment, Aap, which was used to determine the resistance curve

slope, 3J/9a, and the material’s dimensionless tearing modulus, T, where

General Microstructural Assessment

The microstructure of the as-received weld consisted of columnar austenite

grains aligned in the principal direction of heat flow, which was generally

normal to the surface of the weld. Figure 3A shows a cross-sectional view of

the austenite dendrites, and Figure 3B shows a view parallel to the dendrite

columns. No 6-ferrite is present, as verified by electrolytic etching with KOH

and also by magnetometry. However, a solute-redistribution effect occurred

during weld solidification, and this is evident in Figure 3A.

As shown by the unetched sample of Figure MA, a homogeneous distribution of

inclusions was present in the weld. The inclusions have various sizes and are composed mainly of Mn, S, Si, Ti and 0, in various proportions. The average

diameter, d, is 0.64 ym and the density, n, is 9664 inclusions per mm^. In

base metals, the inclusions are typically larger_and fewer in number, as shown

in Figure 4B for AISI 31 OS. For this material, d = 1 0 ym and n = 225/mm^.

The high inclusion density of the weld adversely affects the fracture

properties at 4 K, as described in the discussion.

Mechanical Properties

The 4-K tensile properties of the Fe-l8Cr-l6Ni-6.5Mn-2.4Mo weld deposit

are listed in Table 3. The data scatter for three replicate tests is modest,

and each weld specimen displayed the cup-and-cone morphology that is typical

for ductile fracture in uniaxial tension. The tensile properties of a wrought

austenitic stainless steel base metal, AISI 310S ( Fe-25Cr-20Ni-1 7Mn-0 . 2Mo )

,

are included in Table 3 f or comparison. At 4 K, the average yield strength of

the weld (842 MPa) is slightly hi^er than that of AISI 31 OS (817 MPa);

however, its ultimate strength, elongation, and reduction of area are somewhat

lower

.

Figure 5 presents a typical load- versus-deflection curve for a J-integral

test specimen of the weld at 4 K. The controlled periodic compliance

unloadings and the uncontrolled load-point instabilities give these test

records a serrated appearance. As explained elsewhere [8], serrations are common during the ductile fracture of steels at 4 K and may represent either

unstable plastic deformation alone or unstable plastic deformation combined

with crack extension.

( 2 )

RESULTS

259
Figure 6 presents a typical J-resistance curve for the weld deposit at

^ K. In accordance with standard procedure [6], the regression line is

character ized by a linear least-squares fit and the blunting line is specified

by the expression

J = 2Aaof, (3)

where the flow strength, of, is the average of the uniaxial tensile yield, ay,

and ultimate strengths, ay:

According to ASTM E 8l3“8l, at least four J-versus-Aa points must fall

within the exclusion interval defined by the offset lines at Aa = 0.15 and

1 . 5 mm (see Fig. 6). This requirement was barely satisfied in our tests; no more than four points were obtained because unstable plastic deformation

associated with serrations at 4 K limited the number of unloadings that could

be performed. From the slope of the line connecting and Aaf (the final

physical crack extension increment [?]), the tearing modulus of the weld metal

was determined from Equation (2) to be 50.

The Jjq measurements and Kjq values (estimated from Eq. 1) for each

specimen at 4 K are listed in Table 4. The three results show mild scatter

and an average Kjq(J) value of 190 MPa*m^2. Data for wrought austenitic

stainless steels [6,7] are compared with those of other austenitic stainless

steel welds in Figure 7 [2-4]. In terms of toughness at a given yield

strength, the weld of this study (1) falls about 2Q>% below typical values for

conventional wrought steels, (2) is within the range reported for seven other

fully austenitic stainless steel welds, and (3) clearly surpasses the

toughness of SMA welds containing 5-ferrite.

Fractography

Fracture of the weld at 4 K occurred exclusively by a ductile failure

mechanism, as verified by examination of the fractured J test specimen of

Figure 8. Macroscopically, the 4-K fracture surface is rougher than that of

austenitic stainless steel base metals [6]. Observations at higher

magnifications, as in Figure 9A, show that the fracture surface of the compact

specimen is completely covered by a uniform distribution of very fine dimples,

less than 5 pm in diameter, originating at inclusions in the austenite matrix.

The dimple size is much smaller than typically observed for wrought stainless

steels, such as AISI 31 OS (see Fig. 9B) , where the average dimple is about

35 ym in diameter.

Scanning electron microscope examinations of the weld tensile and compact

specimen fracture surfaces showed no marked difference in the dimple density

or size for fracture in the uniaxial tension or plane strain stress states.

Furthermore, magnetic torsion balance measurements revealed no detectable

magnetic phase on the fracture surfaces. Thus, the Fe-l8Cr-l6Ni-6.5Mn-2.4Mo

weld is completely ductile and fully stable with respect to the Y-to-a’

260
martensitic phase transf ormation at 4 K. There is no evidence of the

"corduroy" fracture mechanism that was recently reported in Mn-modified

metastable type 31 6L weld metal [1].

In Figure 10, mi crostr uctural views of the crack front profile at

midthickness of another J test specimen show that the main crack has deviated

significantly from the horizontal plane, which is the expected plane of crack

propagation. Evidently the crack progresses by successively linking up with

individual voids or void complexes, and the material fractures along strings

of inclusions, not along grain or weld pass interface boundaries or the plane

of maximum stress.

Fatigue Crack Growth Rates

Figure 11 presents the fatigue crack growth rate data for the

Fe-l8Cr-l6Ni-6.5Mn-2.4Mo weld at 4 K. The rates follow a typical power-law

expression , ^ = C(AK)m. (5)

For a logarithmic plot of da/dN versus stress intensity factor range, AK, the

coefficient C is the ordinate intercept at AK = 1 , and m is the slope.

As shown in Figure 11, the weld data extend three orders of magnitude from

da/dN = 10”'^ to 10“^ mm/cycle. A bilinear representation is indicated, but the

data below 10“^ mm/cycle (AK = 28 to 60 MPa*m4) are of primary interest. For

these data, the appropriate power law expression is

^ = 1 .22 X 10~^^(AK)®* °. (6) dN

The fatigue crack growth rates specified by this expression are lower than those

for typical base metals, including AISI 310S and AISI 316 stainless steels. For

example, in Figure 11, the trend line for a representative AISI 316 steel is

given by the expression [9]

da

dN

2.1x1 0"'°( AK) ®. (7)

DISCUSSION

Comparison with Other Austenitic Welds

Data for nine fully austenitic stainless steel welds at 4 K are available

[2,3] and are shown in Figure 7. In general, three of the nine welds match the

base metal strength and toughness levels whereas the rest fall below par by

margins like that reported for the weld tested in this study. Unfortunately,

detailed metallographic and fractographic characterizations were not reported

for the previously tested fully austenitic welds [2,3], and therefore such

evaluations are not possible.

261
Comparison with 316 Welds

The test weld's toughness exceeds the values for conventional type 316

stainless steel welds at comparable strength - levels . The band in Figure 7 for

type 316 compositions represents tests of twenty-three welds produced with

varying ferrite levels (0 to 12 FN) and five different welding processes [4],

The test weld is tougher because it does not have the cleavage- type mechanisms

that cause the 316 weld to be brittle.

Comparison with 31 OS Base Metal

Compared with welds of AISI 31 OS austenitic stainless steels of nearly

equivalent yield strength [8,10], the test weld at A K has a lower ultimate

tensile strength, a lower tensile ductility, a lower Jjc> ^ lower tearing

modulus, and a finer dimple size. Apparently these property differences are due

primarily to the effect of inclusions on the ductile fracture process.

This comparison is meaningful since both the weld metal and the AISI 31 OS steel

are fully austenitic and stable; both materials fail by completely ductile

mechanisms and are free from the effects of martensitic phase transformations,

which would otherwise complicate the interpretation of the fracture process [1 ].

Ductile fracture involves the initiation, growth, and coalescence of voids

to produce a dimpled fracture surface. Typically, voids initiate at Inclusions

at low plastic strains via decohesion at the inclusion/austenite matrix

interfaces. The voids then grow during continued plastic deformation until the

intervening ligaments finally rupture to complete the fracture process. In

wrought stainless steel base metals, such as AISI 31 OS (Fig. AB) , the inclusion

content is relatively low, and the mean free path between Inclusions is

relatively large. Therefore, void growth plays a major role in the fracture

process, and the final dimple size is large (Fig. 9B) . In the weld metal,

however, the inclusion content is relatively hlgi (Fig. 4A); voids are initiated

in large numbers, and they do not grow very large before coalescing (Fig. 9A).

This means that significantly less plastic work is required to accomplish void

coalescence in the weld, and this explains the differences in uniaxial tensile

properties (lower o^, elongation, reduction of area) and fracture toughness

parameters (lower Jjq, lower T).

The estimated volume fraction, fp, of inclusions was calculated in the

following way: We counted the number of inclusions, N, in a 0.1 mm^ area of a polished cross section that were greater than 0.5 pm in size. The average

diameter, d, of the inclusions is the average of ten measurements. The small

number of measurements was felt to be adequate because the inclusions in the

test weld were uniformly small, and in the 31 OS sample, the inclusions were considerably bigger and relatively uniform.

An areal analysis was used for the final calculations of fp:

„ ^

E Ag

P \ ’

where the total area, Aj = 0.1 mm^ and total area of inclusions,

EAa = N TT dV4.

( 8 )

262
For the weld, the estimated volume fraction of inclusions is 0.003 and

the dimple size is about 4 ym in diameter. For the 31 OS wrought steel, fp is

0.0002, and the dimple size is about 35 ym. The flow stress for the two

appears to be similar, about 1000 MPa.

Aside from the inclusion population, the Fe-1 8Cr-1 6Ni-6. 5Mn-2. 4Mo weld

and the AISI 310S base metal have dissimilar grain morphologies. The base

metal grains are equiaxed, whereas the weld grains are columnar. Data for

wrought austenitic stainless steel show that the grain size affects the yield

strength and fracture toughness at 4 K [11]. However, yield strength and

fracture toughness are not independent.

The 4-K fatigue crack growth resistance of the SMA weld is significantly

higher than that of the AISI 316 stainless steel for the specific orientation

and AK range studied (Fig. 11). The SMA weld offers superior performance and

has a higher n value (6.0 versus 3*8) and a lower C value (1.22 x 10“^“ versus

2.1 X 10“^°) than those of AISI 316. The differences in rates for the two

materials range from factors of 1 1 at AK = 28 MPa*m^2 to 1.7 at AK = 60 MPa •m^2 .

In this case, the base metal had a TL orientation and the crack of the

weld was centered in the fusion zone. In both materials, the loading was transverse to the longitudinal plate direction, and crack propagation was in

the longitudinal direction. Previous results [10,12] for austenitic stainless

steel base metals indicated that orientation and compositional variations are

not strong factors if the fatigue crack growth failure modes are ductile. For

example, the measured n values for 14 austenitic stainless steels of the AISI

300 series range from 2.9 to 4.25 at 4 K. Welds, however, appear to be more

sensitive to crack location and specimen orientation. Previously, a high n value (5.3) was reported for an Fe-13Cr-19Mn austenitic weld at 76 K [13]. In

that case, the crack propagation was longitudinal to the fusion line, as in

the present study where n = 6. Apparently the higher n values and superior

fatigue crack growth properties of welds in the longitudinal orientation may

be due to residual stress effects. If so, the performance of welds relative

to base metals should be evaluated thoroughly by considering other specimen

orientations

.

CONCLUSIONS AND RECOMMENDATIONS

The fully austenitic Fe-l8Cr-l6Ni-6.5Mn-2.4Mo SMA weld metal of this

study at 4 K has a yield strength of 842 MPa, a plane strain fracture

toughness Kjq(J) of 190 MPa*m^2, and a tearing modulus of 50. This weld

fails by means of a completely ductile mechanism at 4 K and is more fracture

resistant than type 308L and 31 6L SMA welds containing brittle second-phase

6-ferrite. However, its properties are still lower than those of conventional

wrought base metals. Compared with an AISI 31 OS austenitic stainless steel of

equivalent yield strength, the Fe-l8Cr-l6Ni-6.5Mn-2.4Mo SMA weld at 4 K has a lower ultimate tensile strength, a lower tensile ductility, a lower Jjc, a lower tearing modulus, and a finer dimple size. These differences appear to

be due to the effects of a large inclusion population on the ductile fracture

process of the Fe-l8Cr-l6Ni-6.5Mn-2.4Mo weld.

263
A high inclusion density is typical of SMA welds, especially those with a rutile (type 16) coating. The inclusion density can be reduced by using

shielded metal arc electrodes with calcium oxide/ fluoride (type 15) slag

system or gas metal arc or gas tungsten arc welding. A reduction of the

inclusion size and density is expected to improve the fracture toughness of

the weld deposit without changing the chemical composition of the alloy.

ACKNOWLEDGMENT

The weld was fabricated by Hitachi Research Laboratory, Hitachi, Ltd.,

Hitachi, Japan and provided to NBS through the courtesy of Mr. T. Matsumoto;

we are indebted to him and also to Mr. C. McGowan of NBS who conducted the

tensile tests. The experimental work was supported by the Department of

Energy, Office of Fusion Energy.

REFERENCES

1. J. C. Lippold, M. C. Juhas, and E. N. C. Dalder, The Relationship Between

Microstructure and Fracture of Fully Austentic Type 31 6 L Weldments at

4.2 K, Metall. Trans. A., vol. 16A, (1985), pp . 1835-1848.

2. J. W. Elmer, H. I. McHenry, and T. A. Whipple, Strength and Toughness of

Fully Austenitic Stainless Steel Filler Metals at 4 K, in: Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures - IV,

NBSIR 81-1 645, National Bureau of Standards, Boulder, Colorado (1981),

pp. 289-302.

3. T. A. Whipple, H. I. McHenry, and D. T. Read, Fracture Behavior of

Ferrite-Free Stainless Steel Welds in Liquid Helium, Weld. J. Res.

Suppl., vol. 60, No. 4 (1981), pp . 72s-?8s.

4. R. L. Tobler, T. A. Siewert, and H. I. McHenry, Strength-Toughness

Relationship for Austenitic Structural Steel Welds at 4 K, in: Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures - IX,

National Bureau of Standards, Boulder, Colorado (1986).

5. D. T. Read and R. L. Tobler, Mechanical Property Measurements at Low

Temperatures, in: Advances in Cryogenic Engineering - Materials

,

vol. 28, Plenum Press, New York (1982), pp . 17-28.

6. Standard Test Method for Jjc’ ^ Measure of Fracture Toughness, ASTM

Designation E 813-81, 1984 Annual Book of ASTM Standards, Amer. Soc.

Test. Mater., Philadelphia (1984), pp . 763-781.

7. R. L. Tobler, D. T. Read, and R. P. Reed, Strength-Toughness Relationship

for Interstitially Strengthened AISI 304 Stainless Steels at 4 K

Temperature, in: Fracture Mechanics: Thirteenth Conference, ASTM STP

743 , R. Roberts, Ed., Amer. Soc. Test. Mater., Philadelphia (1981), pp.

250-268.

264
8. R. L. Tobler, Ductile Fracture with Serrations in AISI 31 OS Stainless

Steel at Liquid Helium Temperature, in; Elastic-Plastic Fracture: Second

Symposium, Volume II - Fracture Resistance Curves and Engineering

Applications, ASTM STP 803, C. F. Shih and J. P. Gudas, Eds., Amer. Soc.

Test. Mater., Philadelphia (1 983), PP • 763“776.

9. R. L. Tobler and R. P. Reed, Fatigue Crack Growth Resistance of

Structural Alloys at Cryogenic Temperatures, in: Advances in Cryogenic

Engineering, vol. 24, Plenum Press, New York (1 978), pp . 82-90.

10. W. A. Logsdon, J. M. Wells, and R. Kossowsky, in: Proceedings, Second

International Conference on Mechanical Behavior of Materials, Amer. Soc.

Met., Metals Park, Ohio (1 976 ), pp . 1 283~1289.

11. R. P. Reed, P. T. Purtscher, N. J. Simon, and R. L. Tobler, Alloy 316LN

for Low Temperature Structures: A Summary of Tensile and Fracture Data,

in: Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures— IX, National Bureau of Standards, Boulder, Colorado (1 986 ).

12. R. L. Tobler, H. I. McHenry, and R. P. Reed, Fracture Parameters for an

Iron-1 3? Chromium-19^ Manganese Stainless Steel and Its Welds at

Cryogenic Temperature, in: Advances in Cryogenic Engineering, vol. 24,

Plenum Press, New York (1 978), pp . 560-572.

13. 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. Mater., Philadelphia (1 985), pp . 3~30.

Table 1. Chemical compositions of weld deposit and base metal (wt.5^).

C Si Mn P S Ni Cr Mo N 0

Weld Deposit 0.062 0.29 6.50 0.01 5 0.005 16.30 1 7.64 2.38 0.036 0.042

Base

Material

0.05 0.52 1 .26 0.031 0.006 8.77 18.25 —

265
Table 2, Welding conditions.

Current

(A)

Voltage

(V)

Speed

(cm/min)

Interpass Temperature

(°C)

1 40-1 50 23-24 15 ^150

Table 3. Tensile properties of the weld deposit at 4 K and data for AISI

31 OS for comparison.

Specimen Yield Strength, Ultimate Strength, Elon gation Reduction of Area

No. Oy (MPa) On (MPa) {%) {%)

1 875 1200 34 30

2 824 1 200 NA^ 29 a

3 828 1100 30 22

Average 842 1166 32 27

AISI 310S [ 8 ] 645 1259 36 31 AISI 310S [9] 81 7 1303 64.3 44.6

^broke outside gage marks.

Table 4. Fracture toughness test results for the weld deposit at 4 K.

Specimen No. Thickness

,

B

(mm)

Width,

W

(mm)

Crack Length,

ao^

(mm)

’^Ic (kJ/m^)

Kic(J)

(MPa *m ^/

^

)

Tearing

Modulus

,

T

1 25.425 50.775 31 . 250 1 87 196 —

2 25.450 50.775 30.825 175 1 90 —

3 25.425 50.775 30.900 168 186 50

Average: 177 1 91 —

AISI 310 s [8] — — — 360 285 60

^Computer-calculated crack length.

266
A B

12.7

Tensile

Figure 1 . Specimen geometries used in this study: (A) compact J-integral

test specimen, (B) cylindrical tensile specimen, and (C) fatigue

crack growth rate specimen (all dimensions are in mm).

Figure 2. Specimen orientations with respect to the weld.

267
Figure 3.

100/im

B

62.5 fjim

Weld microstructures, revealed by electrolytic etching in a oxalic acid solution: (A) transverse to the weld direction,

showing cross-sectional view of the austenitic dendrites, and

(B) normal to the weld direction, showing the dendrite arms.

268
Figure 4.

A 50 /jLm

a

B

50 yLim

Inclusions: (A) in the present weld, and (B) in a wrought AISI

310S alloy [?].

269
Deflection, J, mnn

Figure 5. Load-deflection curve for a 25 . 4 -mm- thick compact specimen showing

compliance unloadings and serrations during single-specimen

J-integral tests at 4 K.

270
1000

Figure 6. J-resistance curve for the weld deposit at ^4 K.

271
Yield Strength, Oy, MPa

Figure 7. Comparison of fracture toughness and yield strength properties for

fully austenitic stainless steel welds and base metals at 4 K.

272
Figure 8. Fracture surface appearance of a 25 . 4 -mm-thick compact specimen

tested at 4 K.

273
A

83 /im

B 83^m

Figure 9. SEM view showing dimples on the compact fracture surface of:

(A) Fe-l8Cr-l6Ni-6.5Mn-2.4Mo weld metal, and (B) AISI 310S

base metal [7].

274
B

125 fjLU\

Figure 10. Crack tip profile of a J-integral fracture test specimen showing

voids initiated at inclusions ahead of the advancing crack, at low

and high magnifications, A and B.

275
STRESS INTENSITY FACTOR RANGE. AK, MPa-m^^^

Figure 11. Fatigue crack growth rate data for the SMA weld compared with

typical results for an AISI 316 austenitic stainless steel base

metal

.

276
NONMBTALLiCS

111

NONMETALUCS PROGRAM

LEADER: M. B. Kasen

CONTRIBUTORS: M. W. Austin, R. D. Kriz, H. M. Ledbetter, M. Lei

OBJECTIVES

• Measurement, modeling, and understanding of the properties of composites

that enter essentially into the conception, design, construction,

operation, and performance of magnetic-fusion-energy equipment

• In cooperation with industry, the establishment and refinement of

cryogenic-grade specifications for insulating laminates and test methods

for materials selection

RESEARCH PAST YEAR (1985)

• Finite element analysis to explain the influence of weave geometry in woven

laminates on edge stress singularities

• Determination of Young's modulus and internal friction of a uniaxially

fiber-reinforced composite

• Refinement of the methodology for producing and testing standard composite

specimens to be used in cryogenic neutron irradiation studies

RESEARCH THIS YEAR (1986)

• Study of the relationship between damage and load-deformation behavior of

woven composites at low temperatures for cyclic mechanical and thermal

loads

• Study of the cryogenic neutron-radiation-induced degradation of the

mechanical properties of four organic-matrix composite systems

• Development of a computer program to calculate the load-deformation behavior

of metallic-nonmetallic laminate structures at low temperatures

279
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IsHtJ -

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231

r-1^; .V

'I'

IRRADIATION EFFECTS ON ORGANIC INSULATORS

* +

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The system for producing standard 3.2 mm (0. 125-in) diameter rod

specimens discussed in previous reports^ has been further refined

to permit the fabrication of both f iber-reinforced and neat-resin

specimens from hot-melt resin systems. The method has been suc- cessfully used to produce very high quality specimens duplicating

the G-11CR system and specimens from a variant of that system

eliminating a boron-containing additive. We have also produced

specimens from an epoxy system suitable for impregnation or potting operations and from a bismaleimide polyimide system. These

materials will be used in the first irradiation program in the

National Low Temperature Neutron Irradiation Facility (NLTNIF)

reactor at Oak Ridge. We have refined the 4-K torsional shear

test method for evaluating radiation degradation of the fibermatrix interface and have developed a method of quantitatively

measuring changes in fracture energy as a function of radiation

dose. Cooperative work with laboratories in Japan and England in

this area is continuing and plans are being formulated for joint

production, irradiation, and testing of specimens.

OBJECTIVES

The overall objective of this work is to contribute to development of

organic insulators having the cryogenic neutron irradiation resistance re- quired for MFE systems utilizing superconducting magnet confinement. The

first objective is to develop a component data base from which the optimum

components may be selected on the basis of performance versus cost. The

second objective is to develop a design data base that will provide engineering values of the thermal, electrical, elastic, and mechanical properties of

functional insulating laminates constructed with components selected from the

component data base.

*Work supported in part by the Office of Fusion Energy, U.S. Department of

Energy.

+To be published in "Eighth Annual Progress Report on Special Purpose Material?;

for Magnetically Confined Fusion Reactors," U.S. Department of Energy (1986).

281
PROGRESS AND STATUS

Specimen Developmegt

As described in previous Annual Progress Reports,^ the selected specie- men configuration is a 3.2-mm (0.125-in) diameter rod that can be produced

either as neat (unreinforced) resin or as a uniaxial fiber-reinforced composite. Specimens can be cut to length, irradiated and tested without

further machining.

The quality of specimens currently being produced is illustrated in

Fig. 1 . This quality fully meets the program objective of eliminating

flaws as a complicating factor when evaluating the significance of specific

material parameters on radiation'^induced mechanical property degradation.

We have thus far produced the materials listed in Table 1 .

The G-11CR material duplicates that of the commercial product and is

intended to serve as a reference standard against which the performance of

other materials can be compared. The G‘-11CR(M) modification eliminates the

boron-containing component that serves as a cure accelerator and is ex- pected to improve the neutron irradiation performance of this material.

The Rutherford system was suggested by Dr. David Evans of the Rutherford

Appleton Laboratories in England as a candidate for magnet applications

requiring Impregnation or potting operations. Each of these systems has

been produced as neat^-resin specimens and as specimens reinforced with 48

volume fraction (v/o) of type E glass.

The chemical description of the selected polyimide system is 4,4'-

bismaleimidodiphenylmethane . Prior studies^ have indicated that laminates

fabricated with a bismaleimide polyimide system having the chemical description poly-aminobismaleimide (Keramid 601 ) can provide substantially improved

intrinsic radiation resistance as compared to that of the epoxies, but that

such material has generally inferior mechanical properties —in particular,

very low interlaminar shear strength. We include a different type of polyimide resin in the present program to determine if there exist significant

differences in cryogenic irradiation performance among polyimides differing

in chemical composition.

Specimen production follows the flow diagram given in Fig. 2. Blending

of the G-11CR and G-11CR(M) resin systems was performed ultrasonically to

eliminate the need for solvents and to prevent premature polymerization of

the system. The setup Illustrated in Fig. 3 permits two 406-mm (l6-in)

long reinforced specimens to be produced at the same time. Experience has

shown that the critical operations are thorough degassing, thorough vacuum impregnation of the fibers, and curing under 420 kPa (80 psi) of nitrogen gas

pressure while maintaining a head of resin above the mold. It is also essential that the 3.2-mm (0.125-in) l.d. TFE tubing that forms the mold has a very smooth interior.

282
Test Methodology

It is anticipated that the majority of the radiation-induced damage will occur at the interface between the glass fiber and the resin matrix. Test

methods sensitive to such degradation were therefore emphasized. Candidate

methods are short-beam shear, torsional shear and longitudinal compression.

Development of the short-beam shear test method was described in previous Annual Reports.^ While such a test appears useful, the complex stress

state makes it difficult to quantify results. We therefore concentrated on developing apparatus and techniques for performing torsional shear tests at

cryogenic temperatures . We also developed a test method that provides quantitative data on changes in fracture energy with changing radiation fluence.

A compression test method compatible with the specimen design has already

been developed.^

Torsional shear

This method is attractive because, as applied to a uniaxially reinforced

specimen, the symmetry causes the resolved shear stress at the f iber-matrix'

interface to increase systematically with increasing angular rotation. For a given rotational angle, the magnitude of the shear component will be inversely

proportional to the specimen gage length and proportional to the distance of

the interface from the specimen centerline. As illustrated in Fig. 4(a),

failure initiates at the specimen surface and progresses inward. The interfacial nature of the failure mode is evident in Fig. 4(b).

A typical torsional shear failure at 76 K is shown in Fig. 5(a). In

most cases, seven to ten longitudinal cracks were observed. Some specimens

burst open, revealing fiber bundles and indicating a lower integrity of the

fiber-matrix interface. A typical 76-K load-deflection curve is illustrated

in Fig. 5(b). Because dead weight loading was used, cracking was manifested

as a series of discontinuous rotational displacements.

As illustrated in Fig. 6, the specimen, held in aluminum end caps, is

inserted into a 12.7"mm (0.5~in) diameter internally polished stainless steel

tube and pinned at the bottom. The torsional load is applied by a 6.35“mm

(0.25“in) diameter stainless steel rod which mates with the upper end cap

through a slot. In the developed version of this test method, the load will

be monitored with a torsional load cell, while angular displacement will be

monitored by a geared rotary potentiometer. A motor drive will provide a constant torsional displacement with time.

We investigated the feasibility of freezing the specimens into the end

caps as an alternative to potting with epoxy. This would reduce testing time

and cost and would permit recovery of the undamaged parts of the specimens

for additional testing by other methods. It would also allow specimens to be

periodically removed from the end caps during a test run for quantitative

analysis of damage accumulation by dynamic modulus and damping techniques,

should this be desired. We found that several types of aqueous slurries provided sufficient bond strength when frozen to allow the specimen to be broken, but progressive cracking of the bond within the end caps increased the

283
effective gage section, complicating calculation of the torsional modulus.

We solved this problem by incorporating a series of set screws into the end

caps near the gage section, as illustrated in Fig. 5(a). Contact of these

screws with the specimen prevented rotation and confined the failure to the

gage section. This system will be used in all further torsional tests.

Table 2 summarizes ultimate strength values obtained during ?6“K

torsional testing of neat-resin and reinforced specimens with materials 1 through 4 (Table 1). A somewhat greater coefficient of variability, CV, is

found in neat-resin strength as compared to that of the reinforced materials.

This is expected owing to the low strain to failure of the embrittled resins.

Removal of the BF^MEA component had no significant effect on the strength of

either unreinforced or reinforced 0^*1 ICR materials. The Rutherford resin

developed the highest intrinsic strength at 76 K, but the lower strength of

the reinforced product made with this resin, as compared with that of G-11CR,

suggests a lower fiber-matrix bond integrity. The strength of the blsmaleimide resin was similar to that of G-11CR.

Fracture energy

Although mechanical tests provide useful information, the results are difficult to quantify. We therefore worked to develop a test method that

provides information on the fracture energy, Gj^^, of the reinforced rod-shape

specimens as a function of changes in the specimen condition. The principle

of the method is illustrated in Fig. 7(a). A razor blade is pressed M mm

into the end of the specimen, causing a crack to propagate lengthwise down

the specimen. The fracture energy, Gj^, can then be calculated by the

equation:

Gj^ = 0.01372 d^Ey^/x^ (1)

where E is the Young's modulus, and d, y, and x are as defined on Fig. 7(a).*

Initial test results using the prototype fixture illustrated in Fig. 7(b) are presented in Table 3* All specimens contained 48 v/o of type E glass.

The 76-K toughness of the flexibilized resin specimen (Epon 828/Versamid 140) was found to be much higher than that of the standard G‘-11CR formulation. The

Rutherford potting resin system was also slightly tougher than G-11CR at 76 K. But cooling the Rutherford resin specimen to 76 K decreased its toughness to

about 34)^ of its room temperature value. These changes are in the expected

direction, and the derived Gj^ values are comparable to the 1.9^3. 2 kJ*m ^ values reported for the fracture energy of similar materials tested at room temperature by the double cantilever beam method. The test method therefore

holds promise for quantitatively following the influence of 4-^K irradiation on degradation of the matrix resin and of the f iber'-matrix interface.

Equation (1) is based on the original work of Obreimoff,^ who developed a similar expression for evaluating the splitting strength of mica.

*The author acknowledges the contribution of R. D. Kriz in the derivation

of this equation.

284
Barlow and Windle^ subsequently suggested that the basic method could be

applied to determining the of uniaxially reinforced composite materials. However, their mathematical derivations assumed a geometry in which a thin layer was split off a block of material that remained undeformed. The

derivation of eq. (1) differs in that the moment of inertia was calculated

about the centroidal axis of a semicircular cross section to take into

consideration the geometry of the present specimens. Also, Obreimoff^

defined the fracture energy, Q, as the energy release required to create

one of the two new surfaces. This was equated to Gj^ by Barlow and

Windle,^ and we have followed this convention in deriving eq. (1) to be

consistent with the previous authors. It should be noted, however, that

this definition of fracture energy differs from the current convention in

which the term Gj^ is used to define the energy required to create both

surfaces. Thus, strictly speaking,

Q = Gjq/ 2 = (External Work - Strain Energy)/2.

FISCAL YEAR 1986 PLANS

( 2 )

We plan to cooperate with Dr. R. R. Coltman, Jr., ORNL, in performing a 4.2-K irradiation run in the NLTNIF reactor using the materials listed in

Table 1 . This will involve irradiating both unreinforced and reinforced

specimens at several fluence levels. For this initial experiment, the specimens will be warmed to room temperature after irradiation and will be subsequently cooled to 4.2 K for testing by short-beam shear and torsional shear.

The fracture energy of the reinforced specimens will also be determined. A

comparable series of specimens will be tested in the unirradiated condition.

A total of 320 specimens will be used.

We shall continue cooperative work with the Institute of Science and

Industrial Research, Osaka University, Japan. This group, headed by Prof.

Toichi Okada, includes participants from Kyoto University Research Reactor

Institute, the University of Osaka Prefecture, and the Hoxan Corporation.

This group has developed procedures for producing standard specimens identical to those developed by NBS and is investigating the use of four-point

bending as a test methodology. ° They have investigated the unirradiated

cryogenic performance of specimens reinforced with glass, graphite, silicon

carbide, and aluminum oxide fibers, studying physical as well as mechanical

properties. Their work has shown that aluminum oxide-reinforced epoxy

material has the lowest integrated thermal conductivity of any composite

material from 295 K to 4.2 K."^ These data suggest that such material will

be superior to glass-reinforced composite material for fabrication of efficient thermal standoff members in Dewar construction.

Prof. Okada has agreed to produce a series of specimens in Japan to be

used in a second irradiation series. Following irradiation in the NLTNIF

reactor, these specimens will be sent to NBS for testing and evaluation by

personnel from Prof. Okada’ s laboratory.

Since the required specimens can now be obtained from Japan and possibly from England, we do not anticipate substantial additional specimen

production at NBS. This will permit NBS to concentrate on the materials

testing and analysis portion of the study.

285
REFERENCES 1.

M. B. Kasen, "Materials and Test Methods," in Annual Progress

Report on Special Purpose Materials for Magnetically Confined Fusion Reap- tors," DOE/ER-01

1

3 / 3 , pp. 48-52; May 1984; DOE/ER-01 1 3/4 , pp. 51-57; May

1985.

2. G. F. Hurley and R. R. Coltman, Jr., "Organic Materials for Fusion

Reactor Applications," J. Nucl. Mater. 123 & 124, 1327"1337 (1984).

3 . R. E. Schramm and M. B. Kasen, "Cryogenic Mechanical Properties of

Boron-, Graphite-, and Glass-reinforced Composites," Mater. Sci. Eng. 30,

197-204 (1977).

4. J. W. Obrelmoff, "The Splitting Strength of Mica," Proc. Roy. Soc.

AI 27 , 290-297 (1930).

5. C. Y. Barlow and A. H. Windle, "Razor Blade Test for Composite

Toughness," J. Mater. Sci. Lett. 4, 233^234 (1985).

6. T. Okada, S. Nishijima, H. Yamoaka, K. Miyata, Y. Tsuchida,

K. Kuraoka and S. Namba, "Mechanical Properties of Unidirectionally Reinforced

Materials," presented at ICMC Conference on Materials and Composites at Low

Temperatures, Heidelberg, Germany (1984). To be published in Nonmetallic

Materials and Composites at Low Temperatures - 3 > Plenum Press, New York.

7. M. Takeno, S. Nishijima, T. Okada, K. Fujioka, Y. Tsuchida, and

Y. Kuraoka", Thermal and Mechanical Properties of Advanced Composite Materials at Low Temperatures," presented at CEC/ICMC Conference, Boston,

Mass., August 12-16, 1985. To be published in Advances in Cryogenic Engi - neering - Materials, vol. 32, Plenum Press, New York.

286
Table 1 . Experimental resin materials^

1) G“1 1 CR: Araldite 601 0/Eporal/BF^MEA (100/36/14.2 pbw^)

2) G“11CR(M): Araldite 6010/Eporal (100/36 pbw)

3) Rutherford: Araldite 601 0/XU205/D2000 (100/26/37 pbw)

( potting)

44 ) Polyimide: XU292 developmental bismaleimide

^The use of trade names does not imply recommendation or endorsement by the National Bureau of Standards, nor does it imply that

the materials identified are necessarily the best available for

the purpose.

^Parts by weight

Table 2. Torsional shear data, 76 K (5 specimens of each condition)

Material no. Amount of Ultimate CV

glass (v/o) strength (MPa)

1 48 154.4 0.086

2 48 160.8 0.036

3 48 148.6 0.078

1 none 96.6 0.099

2 none 98.3 0.253

3 none 120.8 0.099

14 none 94.7 0.190

287
Table 3. Initial fracture strength, test results on E-glass“

reinforced specimens (48 v/o)

Resin system T (K) X (mm) Gic ( kJ -mm ^

)

Flexibilized epoxy 76 6.0 2.09

G-1 ICR 76 9.0 0.41

Rutherford epoxy 76 8.5 0.52

Rutherford epoxy 295 6.5 1 .52

Fig. 1 . Transparency of 3 • 2“mm'^diameter; specimen containing

48 volume fraction of glass illustrates freedom from

internal defects. Rutherford resin (see Table 1).

288
ig. 2. Flow diagram for standard specimen production.

289
Fig. 3. Equipment for simultaneously producing two standard specimens.

Illustration shows the setup prior to vacuum transfer of the

degassed resin into the impregnation tubes. The impregnation

tubes and the tube molds are maintained at 70“90°C.

290
Fig. (a): Cross section of a S.S-mm-ciiameter, 48 v/o glass-reinforced

specimen fractured in torsion, illustrating fracture initiation at

the surface and propagation inward. (b): Magnified view, illustrating the interfacial failure mode.

Fig. 5. (a): Torsional failure at 76 K. Set screws prevent extension of

the failure zone beyond the 19-mm gage length. (b): Rotational

displacements on load-displacement curve reflect sequential cracking during dead weight loading.

291
Stainless Steel

Drive Shaft

Fig. 6.

Specimen

Schematic of the torsion test method.

End caps and drive shaft are sized to

permit free rotation at cryogenic

temperatures

.

Stainless Steel

Reaction Tube

Aluminum

End Caps

Reaction Pin

Fig. 7. (a): Schematic of fracture energy test method. Distance x is from

the shoulder of the blade to the end of the crack. Dimension y is

half the blade width. (b): Prototype fixture. Specimen is held in

center-bored aluminum end cap; blade is attached to upper end cap

containing a pin to limit penetration of the specimen to 4 mm. Alignment tube is 9.5~mm i.d. linen phenolic.

292
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1
SCREENING THE PERFORMANCE OF ORGANIC INSULATORS

UNDER CRYOGENIC NEUTRON IRRADIATION^'^

Maurice B. Kasen and Robert B. Stoddard

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Specimens and test procedures are being developed for determining the

significant parameters influencing resistance of organic insulators to neutron irradiation at i| K. The specimens are 3 • 2- mm- diameter rods of exceptionally high quality produced by a method allowing a large number of

experimental variables to be evaluated. Flexural and torsional shear tests

perfomed with these specimens indicate that such tests will be useful in

studying cryogenic neutron irradiation damage to the fiber-matrix interface

Results of 76-K tests on unreinforced and glass -fiber reinforced epoxy and

polyimide materials are presented.

INTRODUCTION

A serious material shortcoming in building superconducting magnets for

magnetic fusion energy (MFE; systems is believed to be the inability of con ventional organic insulators to withstand the level df neutron irradiation

to which they will be exposed in a functioning reactor. It is estimated

that portions of the magnets must be exposed to a fluence level equivalent

to 550 to 1000 MGy (MGy = 10® rads) if excessive shielding costs are to be

avoided.^ Although it is technically feasible to fabricate magnets with

much more radiation resistant inorganic materials, the cost may be prohibitive .

A facility providing irradiation conditions approximating those in an MFE reactor has thus far been unavailable. Fortunately, this will be rec- tified with the coming on line of the National Low Temperature Neutron Irra

diation Facility (NLTNIF) currently under construction at the Oak Ridge

National Laboratory. This facility will provide the desired radiation

spectrum at fluence levels that will allow the required experiments to be

performed in a relatively short period of time. The comparatively large ir

radiation volume and associated test chamber will permit simultaneous irradiation of a large number of specimens as well as in-situ postirradiation

testing at H .2 K without warmup.

^Publication of NBS, not subject to copyright.

''"To be published in Advances in Cryogenic in Cryogenic Engineering--

Materials, volume 32.

293
In the United States, the main research effort involves cooperative re- search between individuals at the National Bureau of Standards, the Oak

Ridge National Laboratory, the Los Alamos National Laboratory, and the

Massachusetts Institute of Technology. Within this group, NBS has the

primary responsibility for providing well-character ized research materials,

conducting pre- and postirradiation mechanical test programs, and determining the influence of irradiation on the material failure mode. Informal but

active cooperation in this research has been established with individuals at

Osaka University in Japan and at the Rutherford Appleton Laboratories in

England.

EXPERIMENTAL PROGRAM

Our primary objective is to determine the significant parameters influencing the degradation of glass-fiber reinforced polymer insulators during

neutron irradiation at 4.2 K. Since many parameters must be evaluated, it

is necessary to establish an efficient, systematic screening program.

Specimen Selection and Preparation

Differential radiation heating of the specimen cross section, the need

to minimize residual radioactivity, and the many variables that must be

studied dictate the need for small, high-quality, well-character ized specimens. Surface defects introduced by machining should be eliminated. The same specimen configuration must be available in a variety of unreinforced

(neat) resin specimens and as specimens reinforced with any desired type and

amount of glass fiber. The specimen should lend itself to testing in a variety of modes. There is no commercial source of materials meeting these

requirements. We have therefore developed in-house capability of producing

the required specimens.

We have selected the rod-shaped specimens illustrated in Fig. 1(a) as best meeting these requirements. Glass-fiber bundles are vacuum impregnated

with well-degassed resin pulled into a 300-mm long, 3.2-mm i.d. smooth-bore

TFE mold and cured under approximately 420-kPa nitrogen gas pressure. Neat

resin specimens are produced by vacuum drawing degassed resin into the same molds. As illustrated in Fig. 1(b), specimen quality is very high. Fiber

volume fractions can be calculated from knowledge of the tube dimensions and

glass density. Resin chemistry, cure state, fiber chemistry, and fiber finish can be altered as desired for a given experiment. The specimen surface

finish is excellent, and the coefficient of variability in diameter among

rods produced with a given mold is on the order of 0.001 to 0.003, indicating excellent dimensional reproducibility. These specimens can be cut to

length and directly tested.

Some polymer systems of interest require blending of a liquid resin

with a solid cure agent. In industrial practice, this is facilitated by the

addition of solvents, but this cannot be done with the NBS production method

because evolution of the solvents would cause porosity during curing in the

closed mold. We avoid this problem by the ultrasonic blending of such systems . ^

294
This method was used to produce unreinforced and type E glass-fiber

inforced specimens with the resin systems listed in Table 1. Materials 1 through ^ will be used in the initial radiation program to be conducted in

the NLTNIF reactor at Oak Ridge. The G-11CR system (material 1) provides a baseline of material performance. However, this basic system contains

BF^MEA as an accelerator to shorten the cure cycle. Since boron has a high

neutron capture cross section, this addition may be contributing to excessive degradation of the resin. We therefore produced a variant of the

G-11CR formulation without this addition to evaluate this parameter (material 2). An impregnation resin (material 3) was included in response to suggestions made at the 1984 NBS/DOE workshop that such a system would be of

value to magnet fabricators. The proprietary bismaleimide polyimide system

(material 4) represents the state-of-the-art in advanced resin design for

this type of material. The bisphenol A/polyamide system (material 5) is a simple system used during the development stage of the program and will not

be included in the radiation study.

Test Method Development

Evidence suggests that the primary cause of degradation in the mechanical properties of reinforced insulators will be radiation-induced damage to

the interface between the fiber and the polymer matrix. 3 We therefore

concentrate on developing test methods that are sensitive to such damage.

Primary candidates are short-beam shear and torsional shear tests.

The fixture used for short-beam shear tests is similar to a conventional three-point fixture for apparent interlaminar shear strength tests,

differing primarily in contouring the supports and the loading nose to

accommodate round specimens. At the proper span/depth ratio, this test produces an essentially mode-II type loading in which the shear stress is

resolved parallel to the fiber direction. Required specimen length is only

15 mm for a span/depth ratio of 3.0, making it very economical of material.

Typical load-time curves produced at 76 K and 295 K by this test method are shown in Fig. 2. The load drops appearing on the 76-K curve reflect the

onset of consecutive shear failures with increasing stress. Continuous bending over of the 295 K curve is attributed to the high room temperature

ductility of the bisphenol A/polyimide resin system.

A typical 76-K shear failure is illustrated in Fig. 3. Narrow shear

bands initially progressed into the specimen in an inverted V configuration

from the point of contact with the loading nose. These bands subsequently

acted as nuclei for massive shear failures that propagated along the specimen length, giving rise to the observed load drops.

Although the flexural shear test appears potentially useful, the change

in specimen shape during deformation introduces failure modes which are a combination of tension, compression and shear. This makes it difficult to

quantify the results. An alternative is to conduct tests in torsion. Such a test in a uniaxially reinforced composite also produces a combined shear,

tension and compression loading on the interface, but in this case the

symmetry of deformation simplifies calculation of the various stress

gradients and permits the shear component of the stress to be isolated.

295
The appearance of 76-K failures produced by applying dead-weight torsional loading to unreinforced and to glass-fiber reinforced specimens

bonded into aluminum end caps is illustrated in Fig. 4(a-b). The unreinforced specimens developed typical 45° shear failures. In most cases, seven

to ten longitudinal cracks were observed in the reinforced specimens. But

some specimens burst open, revealing individual fiber bundles. Indicating a lower integrity of the fiber-matrix interface. A typical load-deflection

curve obtained in torsion at 76 K is shown in Fig. 5. Under dead weight

loading, cracks are manifested as sudden jumps in displacement.

We investigated the feasibility of freezing the specimens into the end

caps as an alternative to potting with epoxy. This would permit recovery of

the undamaged parts of the specimens for additional testing by other methods.

It would also allow specimens to be periodically removed from the end caps

during a test run for quantitative analysis of damage accumulation by

dynamic modulus and damping techniques. Although this approach was found to

be feasible using several types of aqueous slurries, it was also found that

specimens prepared in this way might develop progressive cracking of the

bond between the specimen and the end cap from the gage section into the

gripped portion as the torsional load increased. Some of the displacement

jumps in Fig. 5 may reflect such cracking. This approach to the test method

is undergoing further development.

Preliminary Test Results

A series of flexural shear tests was conducted with the bisphenol

A/polyamide material (material 5) at 76 K. Five specimens each from two

rods containing either 48 or 58 v/o type E glass were tested at each temperature at a span/depth ratio of 3*0. The data, summarized in Table 2, re- flect the stress required to initiate the first massive longitudinal shear

failure (first load drop). The low coefficients of variability indicate

good consistency from specimen to specimen and between specimens from different rods. No significant difference in shear strength was noted between

the rods of differing fiber content. It is probable that this reflects a maximum in shear strength in the vicinity of 50 to 60 v/o fiber content. At

very high fiber content there is insufficient resin to ensure fiber separation.

Torsional tests were conducted at 76 K with materials 1 through 4 (Table 1). The ultimate strength values are summarized in Table 3. A somewhat greater variability was found in the neat resin strength as compared

with that of the reinforced materials. This is expected owing to the low

strain to failure of the embrittled resins. Variability among the reinforced materials was slightly higher than for the flexural tests, but still

reasonable. Removal of the BF^MEA component had no significant influence on

the strength of either the unreinforced or reinforced G-11CR materials. The

Rutherford resin developed the highest intrinsic strength at 76 K, but the

lower strength of the reinforced product made with this resin, as compared

with that of G-11CR, indicates a lower fiber-matrix bond integrity. The

strength of the polyimide resin vas about that of G-11CR.

296
These tests appear to be useful in characterizing cryogenic neutron

irradiation damage to the various types of materials that will be evaluated

during this program.

ACKNOWLEDGMENTS

This work is supported by the Office of Fusion Energy, U.S. Department

of Energy, Dr. Marvin Cohen, Project Manager.

REFERENCES

1. 0. F. Hurley and R. R. Coltman, Organic materials for fusion reactor

applications, J . Nucl . Mater . 123 & 1 2^: 1 327“! 337 (1984).

2. W. T. Hodges and T. L. St. Clair, "Ultrasonic Mixing of Epoxy Curing

Agents," Technical Memorandum 85643, NASA Langley Research Center,

Hampton, Virginia (1983).

3. S. Egusa, M. A. Kirk, R. C. Birtcher, M. Hagiwara, and S. Kawanishi,

Irradiation effects on the mechanical properties of composite organic

insulators, Nucl. Phys . Methods Phys

.

B229:6l0-6l6 (1984).

4. A. Udagawa, S. Kawanishi, S. Egusa, and M. Hagiwara, Radiation induced

debonding of matrix-interface in organic composite materials, J . Mat . Sci . Lett. 3:68-70 (1984).

Table 1. Experimental Resin Materials

1 ) standard G-llCR epoxy system

2) Standard G-llCR epoxy without BF^MEA accelerator

3) Rutherford bisphenol A epoxy impregnating resin

4) Proprietary bismaleimide polyimide

5) Bisphenol A epoxy with polyamide cure

Table 2. Flexural Shear Data, Material 5, at 76 K (10 Specimens of Each

Condition)

V/0 Glass Rod No. Ultimate

Strength (MPa) CV

48 1 23.6 0.042

2 23.0 0.087

23.3 0.064

58 1 23.3 0.047

2 23.6 0.013

23.5 0.034

297
Table 3- Torsional Shear Data, 76 K (5 Specimens of Each Condition)

Material No. V/O Glass

Ultimate

Strength (MPa) CV

1 58 158.iJ 0.086

2 58 160.8 0.036

3 58 148.6 0.078

1 none 96.6 0.099

2 none 93.3 0.253

3 none 120.8 0.099

none 94.7 0.190

Fig. 1. (a) Typical 3 • 2-mm-d iameter roc specimens. (b) An indication of

the rod quality as revealed by its transparency . Top to bottom in

(b) are neat polyimide, reinforced G-11CR, neat G-11CR, reinforced

Rutherford resin, and neat Rutherford resin. Reinforced specimens

have 58 v/o glass content.

298
Fig. 2. Typical lead-time curves obtained during

flexural shear tests. Serrations at 76 K

indicate onset of longitudinal shear.

fig. 3. Typical longitudinal shear failure at 76 K. Arrow indicates the

point of contact of the loading nose. Shear bands have propagated

into the specimen, forming an inverted V and acting as the

nucleating point for massive longitudinal shear failure.

299
Fig. (a) Typical 76-K torsional failure in an unreinforced resin

specimen illustrating the method of holding in aluminum end caps,

(b) Typical 76-K torsional failures in glass-reinforced Rutherford

resin system (material 3, Table 1).

Fig. 5. Torsional load-deflection curve obtained at

76 K illustrating sequential interfacial

crack development (A-E).

300
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Young modulus and internal friction of a fiber-reinforced composite

H. M. Ledbetter, Ming Lei, and M. W. Austin

Fracture and Deformation Division, Institutefor Materials Science and Engineering,

National Bureau ofStandards, Boulder, Colorado 80303

(Received 28 August 1985; accepted for publication 3 December 1985)

By a kilohertz-frequency resonance method we determined the Young modulus and internal

friction of a uniaxially fiber-reinforced composite. The composite comprised glass fibers in an

epoxy-resin matrix. We studied three fiber contents: 0, 41, and 49 vol %. The Young modulus fit a

linear rule of mixture. The internal friction fit a classical free-damped-oscillator model where one

assumes a linear rule of mixture for three quantities: mass, force constant, and mechanicalresistance constant.

I. INTRODUCTION

Glass-fiber epoxy-matrix composites enjoy many uses because they possess special properties. One can tailor many

of the physical and mechanical properties for a particular

application by controlling fiber orientation and fiber volume

fraction. As composites increasingly replace conventional

materials, we see strong needs to optimize composite properties. This means understanding the relationships between

macroscopic properties and microscopic properties. To use composites effectively, we must explain and predict their

properties. Because composites show such complexity, so many variables, usual empirical methods fail to produce the necessary understanding. We believe such understanding

will arise only through careful property measurements combined with careful modeling.

The present study considers measurements and modeling of some physical properties of an eposy resin reinforced

with continuous uniaxial glass fibers. We considered three

glass-fiber volume fractions: 0, 0.41, and 0.49. A physical

property of a composite material depends mainly on three

ingredients: matrix property, inclusion property, and phase

geometry. Phase geometry includes many variables: volume

fraction, inclusion shape, inclusion orientation, inclusion

size, and inclusion distribution. For some composites, especially at higher temperatures, the inclusion-matrix interface

affects a physical property.

The present study considers two strongly related physical properties: elastic constants and internal friction. Especially, we focus on Young modulus E and internal friction Q

~ ‘ determined in rod-shaped specimens in a Young-modulus (extensional-wave) mode. E and Q~^ relate to the real and imaginary parts of the elastic-stiffness tensor.

Elastic constants enter many aspects of composite-material behavior: stiffness-weight ratio, load-deflection, elastic

instability, thermoelastic stress, residual stress, sound-wave

velocities, material characterization, relationship to other

physical properties (for example, thermal expansivity and

specific heat), plastic deformation, theoretical strength, and

nondestructive evaluation.

Internal friction of composites relates to many of the

above phenomena. In addition, it relates especially to structural damping and to detection of premonitory failure as manifested in cracking, debonding, and delamination.

1972 J. Appl. Phys. 59 (6), 1 5 March 1 986

Besides measuring these two physical properties, we

modeled their depencencies on volume fraction. For the

Young modulus, especially the Young modulus in the fiber direction E^, a plethora of good models and model-measurement comparisons exist. E^ represents probably the best un- derstood composite-material elastic constant. It yields easily to various measurement methods, and it usually follows

closely a linear rule of mixture. Not so for internal friction Q~^. Although considered extensively from an engineeringdesign viewpoint, the systematics ofQ~^ for composites re- main uncertain. Here we consider principally how £"3 and

the associated Q

~ ’ depend on volume fraction of reinforcing

fiber.

II. EXPERIMENT

A. Materials

Composites were made at NBS using vacuum-impregnation methods described by Kasen.' The manufacturing

method yielded 3-mm-diam cylinders up to 30 cm long. Except for cutting to length, we measured these as-received. The matrix material consisted of a commercial epoxide resin formulated for use in a low-temperature radiation environ- ment. The reinforcing phase consisted of tows ofcommercial

glass fibers 7 //m in diameter with a reported mass density of

2.60 g/cm^ and Young modulus of 72.4 GPa. Figure 1 shows a photomicrograph ofthe 41% glass composite. Table

I contains the mass densities of all three studied materials. Fiber volume fraction was determined from three parameters: glass mass, glass mass density, and mold volume.

B. Measurements

Determining Young modulus E required measuring the sound velocity v and the mass density p, and using the rela- tionship

E = pv^=pf\k^, (

1 ) where f^ denotes specimen resonance frequency and k de- notes resonance wave length. At half-wave resonance. k = 2/j (where 4 denotes specimen length) and

E = ^pf]l], (2.

A Marx^’^ oscillator yielded the experimental resonance fre- quency according to

ifl/?

301
FIG. 1. Photomicrograph of studied composite transverse to uniaxial glassfiber bundles in an epoxy matrix. Fiber diameter equals 7 /rm. Nonhomogeous fiber distribution results from fiber bundles used in manufacture.

= (3)

i where m denotes mass, o denotes the complete oscillator,

and i denotes oscillator components: quartz-crystal driver,

quartz-crystal gage, and specimen. It follows simply from

Eq. (3) that

fl =fl + (/o , (4)

where subscript q denotes quartz crystals. The quartz-crystal length fixes . A maximum voltage Vg , corresponding to a maximum strain amplitude at the gage-crystal midpoint,

detects/,. Sweeping frequency reveals the voltage maxi- mum. Extensional elastic waves were launched along the rod

axis by using — 18.5° Z-cut quartz crystals.

For a sharp, Lorentzian resonance peak the composite

oscillator’s internal friction ^

‘ is

Co”'

(/2 -/,)//.. (5)

the resonance-peak width at \/yj2 of maximum in a strain- versus-frequency plot where strain is proportional to Vg.

Equation ( 5 ) holds for Q

~

' up to 0. 1 or 0.2. Experimentally one observes M*, the modulus in phase with strain. But one

TABLE I. Measured properties of glass/epoxy composites.

c

E, (GPa)

Q-'

(10-’) P

(g/cm’)

L

(kHz)

1. (cm)

0.0 3.114 75.6 1.115 30.47 2.694

0.41 32.18 6.15 1.654 89.03 2.477

0.49 39.12 4.57 1.852 99.94 2.294

usually seeks , the usually more-relevant storage modulus. For zero internal friction, — M*.

Similar to Eq. (3), = . ( 6 )

i Thus, a relationship similar to Eq. (4) yields the specimen’s

internal friction: = CcT'

(0<r' - • (7)

The above relationships apply only near resonance, when

fs=fq=fo- ( 8 ) Figures 2 and 3 show the errors arising from departures from

Eq. (8), expressed as a departure from /,// = 1. In Fig. 2, the Young-modulus measurements are fit to E = a bx + cx^

,

(9)

where

X=fg/fs- ( 10 ) In Fig. 3, we used a similar relationship for Q~'. For both

physical properties and 0

' we took the maximum or minimum value, not the values at /,// = 1. For these materials, these two values differed approximately 0. 1 % for E^

and 1% for 2

We measured mass density by Archimedes’s method us- ing distilled water as a density standard.

Concerning errors, based on the author’s experience,

they estimate the E uncertainty as 1 % and the Q

” ' uncertainty as 5%.

III. RESULTS

Table I contains the study’s principal measurement results: values of E^ and ^

‘ for the three studied materials. The table also contains mass densities together with other

measurement results. Figure 4 shows graphs of E and Q

~ ‘

versus fiber volume fraction c.

IV. DISCUSSION

In this section, we want to consider principally the ex- planation for the dependence of E and Q on volume per- cent inclusion, or fiber. Usually, this requires two models

fq/fs

FIG. 2. Graph of apparent Young modulus vs /,//,. We took optimum

value, rather than /,//, value, as correct.

1973 J. Appl. Phys., Vol. 59, No. 6, 15 March 1986 302

Ledbetter, Lei, and Austin 1973
FIG. 3. Similar to Fig. 2, except for internal friction.

because E reflects an elastic constant or “spring constant” in phase with strain, but reflects the phase difference

between stress and strain. Thus, E and Q represent coupled but disparate properties. One way to express the interre- lationship is ^ = y/y, = , (11)

where y denotes (complex) elastic compliance, M denotes

(complex) elastic stiffness, component 1 is in phase with

stress, and component 2 is out of phase with stress. The real parts of the elastic moduli J^ and A/, represent elastic-ener- gy storage. The imaginary parts J2 and M2 represent elastic- energy dissipation.

First, we consider a model for E^. Using energy methods, Hashin and Rosen^ derived relationships for the effective moduli of a continuous-fiber-reinforced composite

where the fibers are distributed randomly and homogeneously. Hashin^ gave these relationships in essentially the

form

c{kf-k^){k„ +/z„)

H ( 12 )

CO CL<3

liT w

D

_lD

O

O

V)b

z3

O

>

FIG. 4. Young modulus along fiber direction £3 and internal friction ^

~ ‘

vs fiber volume fraction c. Straight line assumes linear rule of mixture and

represents least-squares fit to three measurements. Curved line represents Eq. (30) based on a free-damped-oscillator model and linear rules of mixture for mass, force constant, and resistance constant.

and

E'lt — H r • -ciEf-Em)

( 13 )

Here, k denotes the two-dimensional, plane-strain bulk modulus, which is /I -f /2 in an isotropic material, where A and

// denote Lame constants. Subscript T denotes deformation

in the x, — X2 plane, perpendicular to the fibers. Subscript

L T denotes shear deformation in any plane containing thexj

axis, that is, the fiber axis. Equations (12) and (13) were

also derived by HilL for a single fiber placed concentrically

in a cylindrical matrix. Hermans® obtained them also by considering Hill’s assembly to be embedded in an unbounded

solid having the effective elastic moduli of the composite. In

this model, the fiber : matrix cylinder radius ratio is /',/r2 = c'^^. Bose and MaP’^° also obtained Eqs. ( 12) and (13),

but by a different approach: by considering the effective velocity for long-wavelength plane waves propagating perpendicular to the fibers, for both a longitudinal wave and for a shear wave polarized along x^, the fiber axis. Knowing k-p, one can derive the Young modulus Ep

( along X3) and the Poisson ratio = — 6,763, where uniaxial stress is along X3, and 6 denotes strain. One does this by

using the relationship given by HilL:

{\/kp)-{\/kA {\/kp) - {\/k^)

''LT

{\/kf) - {\/k„)

Vf-V„

~ ~ C{Ej- — E^ ) (14)

Solving for Ep using Eq. (12),

Ep=E2={l-c)E^+cE^

4c( 1 - c)(vy - v„ y

(1 - c)/k^ -F c/k„ + \/n^

(15)

This relationship represents a simple, linear “rule-of-mixtures” plus a higher-order correction term. Using elastic constants given in Table II, we see that, for c = 0.49, the higher-order term in Eq. (15) equals only 0.7% of the rule- of-mixture term. Thus, we neglect it. The straight line in Fig.

4 represents a linear least-squares fit to our three measurements. By extrapolation it gives Ej- = 15.1 GPa, very close to reported values of the fiber Young modulus. [Using Eq.

(15) inversely provides a good way to determine Ef .

)

Second we consider the dependence of the internal fric- tion on the fiber volume fraction. We begin with an equation

for free vibrations":

TABLE II. Properties of constituent materials.

p E n B k (g/cm^) (GPa) (GPa) (GPa) (GPa) u,l>.i)

Glass 2.577 75.7 30.3 50.5 60 6 'O ; n Resin 1.155 3.11 1.17 3.1 1 3 50 : )4 uMt

1 974 J. Appl. Phys., Vol. 59, No. 6, 1 5 March 1 986 Ledbetter, <.ei, and Austr n -s

303
mx bx kx = Q , (16)

where m denotes mass, b denotes mechanical-resistance constant, k denotes force constant, x denotes displacement, and

dots signify differentiation with respect to time. Recognizing

that k relates to the fundamental resonance frequency coq according to o)Q=k /m (17)

and defining

2y = b /m , (18) we rewrite Eq. (16) as X + 2yx + co^x = 0 . (19)

If 0)1 > y^, then Eq. (19) has a solution"

X = C exp{ — yt)sin{o)t + (f)) , (20)

where the frequency co relates to the natural (undamped)

frequency cOq according to

I

0)^ ^0)1- f (21)

and (j) denotes the initial phase. Following French, ' ‘ neglecting the difference between o) and coq, the internal friction is Q

^ = 2y/(Oo = b /{km)''^ . (22)

Following Igata and Kohyama,

of mixture for three quantities:

" we adopt linear rules

m = (1 — c)m^ + ctUf , (23)

A:= (1 —c)k^ +ckf, (24)

and

6 = (1 - c)b„ + cbf , (25)

where subscripts m and / denote matrix and fiber. If the

interface contributes to the damping, it does so through b\ one would add a 6, term to the right-hand side of Eq. (25).

Here, subscript / denotes interface. Combining Eqs. (22)-(25) leads to

Q- _C

1-c b„)

k,\

( " 1

2 ntf kf1 -cl kj ll -cJ

(26)

(27)

where

Recognizing that m =^pV=m„ -\- rrif = cpfV -y - c)p^V

,

(28)

where p denotes mass density and V denotes volume, then

mf Pf

p„

and we rewrite Eq. (26) as

(29)

V \-c b„) V V

\-c k„

\ \ -c/ \ \ - cj p^ k„ j

- 1/2

. (30)

Because the glass fiber is much stiffer and less viscous than

the resin matrix, we make the approximation that bf = 0. Then the second term in the left-most brackets on the righthand side of Eq. (30) disappears. The curve in Fig. 4 represents Eq. (30) with bf = 0. For the force constants, we used

the relationship

k = Ea , (31)

where a denotes a characteristic interatomic spacing. This

led to kf — 24.44 N/m and k^ = 0.31 14 N/m. We treated

as the single adjustable fitting parameter in Eq. (30),

using a„ = I A. Altogether, we find the measurement-model fit for g

“ ’ in Fig. 4 better than expected.

We see no need to postulate a matrix-fiber-interface in- ternal-friction term. In Fig. 4, an interface term would raise the model curve above observation. Through the single adjustable parameter a„ , one cannot lower the curve because

becomes physically unreal.

Finally, we note that the nonhomogeneous fiber distri- bution shown in Fig. 1 does not affect the rule-of-mixture

calculations for the Young modulus in the fiber direction. Elastic constants that follow a rule of mixture do not depend on fiber-distribution homogeneity. Elastic stiffening in fiber- rich regions is compensated exactly by elastic softening in fiber-poor regions. Model calculations by Datta, Ledbetter,

and Kriz" show that other elastic constants depend on fiber- distribution nonhomogeneity. For example, the shear modulus Slid the Poisson ration depend strongly on de- partures from homogeneity.

V. CONCLUSIONS

From this study, we reach four conclusions concerning

a composite material consisting of an epoxy resin reinforced

with continuous parallel glass fibers:

( 1 ) As expected from many previous studies, the fiber- direction Young modulus follows a linear rule-of-mixture.(2) If one does not know the fiber Young modulus Ef, one can determine it reliably by measuring E^ for a few composites with different c values and extrapolating linearly to c= 1.0.

( 3 ) The internal friction Q

~ ‘ follows a classical free- damped-oscillator model that assumes a linear rule of mixture for three quantities: mass, force constant, and mechanical-resistance constant. (4) No evidence emerged that the fiber-matrix interface contributed to internal friction.

ACKNOWLEDGMENT

Dr. M. B. Kasen of NBS produced the composite materials. Owens-Corning (Granville, Ohio) furnished the glass

1975 J. Appl. Phys., Vol. 59, No. 6, 15 March 1986 1975 304

Ledbetter, Lei, and Austin
libers M Lei, a guest worker at NBS, is on leave from Insti- tute of Metals Research, Sheyang, People’s Republic of China.

'M It Kasen, in National Bureau of Standards Report NBSIR 85-3075

( May l<)85). pp. 87-97.

J Marn, Rev. Sci. Instrum. 22, 503 (1951). 'W H Robinson and A. Edgar, IEEE Trans. Sonics Ultrason. SU-21, 8

( 1974). *A S Nowick and B. S. Berry, Anelastic Relaxation in Crystalline Solids (Academic, New York, 1972).

’Z. Hashin and B. W. Rosen, J. Appl. Mech. 31, 223 ( 1964).

'’Z. Hashin, in Mechanics of Composite Materials (Academic, New York,

1970), pp. 201-240.

’R. Hill, J. Mech. Phys. Solids 12, 199 ( 1964).

*J. J. Hermans, Proc. K. Ned. Akad. Wet. B 70, 1 ( 1967).

•"S. K. Bose and A. K. Mai, Int. J. Solids Struct. 9, 1075 ( 1973).

‘°S. K. Bose and A. K. Mai, J. Mech. Phys. Solids 22, 217 ( 1974). “a. P. French, Vibrations and Waves (Norton, New York, 1971), pp. 62-

67. '^N. Igata and A. Kohyama, Trans. Jpn. Soc. Compos. Mater. 2, 30 ( 1976).

'’S. K- Datta, H. M. Ledbetter, and R. D. Kriz, Int. J. Solids Struct. 20, 429

(1984).

»

A

1976 J. Appl. Phys., Vol. 59, No. 6, 15 March 1986 305

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/«
EDGE STRESSES IN WOVEN LAMINATES

AT LOW TEMPERATURES

Ronald D. Kriz

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado 80303

Woven glass-epoxy laminates are used as nonmetallic components at low

temperatures in magnetic fusion energy structures. Previous damage studies

on G-10CR and G-11CR cryogenic grade woven laminates revealed that most of

the damage occurred in the laminated interior. An existing, generalized

plane strain, finite element model was modified to predict stress states at

the laminate interior and free edges. Finite element results demonstrated

that the weave geometry reduces edge stresses at low temperatures. Delamination edge stresses in woven laminates are more sensitive to small changes

in temperature than those in nonwoven laminates. Results suggest that small

warp angles suppress edge damage better than large angles.

Introduction

Woven glass-epoxy laminates are used at low temperatures for a variety

of applications in magnetic fusion energy (MFE) structures. These woven

laminates are used mainly as electrical and thermal barriers that provide

minimal structural support. However, when cooled to liquid helium temperatures (4 K), large internal stresses arise from the differential thermal

contractions of dissimilar components within the laminate. When these

thermally induced stresses are combined with stresses induced by small

mechanical loads (magnetic loads), microcracks may initiate at loads well

below yield. The formation of large numbers of microcracks could limit the

usefulness of these materials for MFE applications.

Previous studies [1,2,3] measured the mechanical properties of two

popular glass-epoxy woven materials, G-IOCR and G-11CR. Later studies [4,5]

related individual damage events to the mechanical performance of G-IOCR and

G-11CR at small strains and low temperatures. All tes-ts were performed on thin laminates where internal and edge damage could be monitored periodically

with increasing load. Results from these investigations showed that bimodulus behavior resulted at strains well below the ultimate strain and that

damage originated in the interior but not at the edge. For nonwoven

laminates, damage in the form of ply cracks and delaminations originates dt

the edge because a large stress state exists there; this edge damage is

representative of damage in the interior. For woven laminates, edge damagis not representative of internal damage.

307
In this study, we modeled edge stresses for woven and nonwoven thin

laminates and determined the influence of weave geometry and thermal effects on edge stresses. With these results, we explain why edge damage differs for woven and nonwoven laminates.

Edge Stress Analysis

Figure 1 defines simple geometries for woven and nonwoven laminates

loaded in tension in the x direction. The symmetric stacking sequence

[0/90]g for the nonwoven laminate is denoted by a subscript (s), and the

ordering of the plies is 0° near the top surface and 90° near the midplane.

Similar stacking sequences exist near the free edge of woven laminates, where

we use a superscript (0) to denote the weave curvature in the y-z plane. Stresses

in the y-z plane develop from the mismatch in Poisson’s contraction between

dissimilar ply or fiber bundle orientations. For a [0/90]^ nonwoven laminate, the edge stress, o^, develops because the transverse contraction in

the 0° ply is greater than the transverse contraction in the 90° ply. Hence,

along the midplane nearest the edge, is in tension because the 0° and 90°

plies are bonded together. To satisfy equilibrium, o must change from

tension to compression as we move along the midplane tz = 0) away from the edge

(y = -B). For the [90/0]^ laminate, the opposite is true: is compressive

along the midplane nearest the edge. The contribution of weave geometry and

thermal loads to the edge stress problem is not easily understood by simple

intuitive models. For a better understanding of the influence of weave geometry and thermal loads on edge stresses, we extended the existing edge

stress model of Pipes and Pagano [6] for some simple weave geometries.

The edge stress problem was first defined by Pipes and Pagano [6] for

nonwoven thin laminates loaded in tension normal to the y-z plane; it is the

generalized plane strain problem. For woven laminates, Kriz [4,5] modified

the generalized plane strain problem to include in-plane loads and fiber axis

rotation, 9, in the x-z plane far from the edge at y

= 0. In-plane and

out-of-plane loads were calculated by simple two-dimensional plate theory.

For woven laminates, the formulation for the edge stress problem in the y-z

plane is identical to the generalized plane strain problem defined in

reference 5, except for a 90° rotation about the z axis. The finite-element

formulation used in this study is given in Appendix A of reference 5. The

only exception to the previous study is that there are no applied loads in

the y-z plane. The same elastic-thermal properties from reference 5 are used

in this finite element study; they are listed in Table 1

.

The finite element grids for the woven and nonwoven edge stress problem

are shown in Fig. 2. For the [90/0]^ nonwoven grid, fibers are oriented at

either 0° or 90° to the load axis (x direction). For the [90/0]| woven

grid, fiber bundles are oriented at either 0° or 90° to the load axis, and

the fiber bundles oriented at 90° can have an additional warp angle, 0, in

the y-z plane.

Both grids shown in Fig. 2 represent only one-quarter of the total

symmetric grid geometry. There are four possible symmetric geometries. For

the nonwoven laminate we can have the sequence [90/0]„, shown in Fig. 2, or O

308
the sequence [0/90]^, which can be obtained by interchanging the top surface

and midplane. Similarly, we classify woven laminate geometries at the edge

as [90/0]® and [0/90]|. Figure 3 shows the total symmetry for these two weave geometries.

Both grids have a higher density of elements near the edge, where we predict a large gradient of stress as y approaches B. Each element is a constant strain triangle with three degrees of freedom per node. The non- woven grid has 2044 elements with 1115 nodes; the woven grid has 2394 elements with 1286 nodes. The global stiffness matrix of the woven grid is

larger, with 3805 degrees of freedom and a half bandwidth of 346.

Both grids model out-of-plane displacements but only for unit depth.

Hence, no out-of-plane curvatures are modeled in this study. We assume that

this approximation is a serious limitation only for the larger warp angles.

Obviously, a three-dimensional finite element model would be more accurate,

but we considered only the influence of weave curvature in the y-z plane on

the edge stresses in this study.

Mechanical and thermal loads were applied to woven grids with 15°, 30°,

and 45° of curvature by increasing the aspect ratio (B/H) of the grid.

Typically, the edge stresses are normalized by the mechanical load (outof-plane strain) , but in this study we superposed mechanical loads on thermal

loads. Hence, we chose a one-percent strain because initial damage was observed at this strain, and we chose a thermal load of 100 K because at this

temperature the in-plane stresses are of the same order of magnitude as those

produced by the mechanical load.

Results and Discussion

In Fig. 4, we compare the stress distribution of the [0/90]g nonwoven

grid in Fig. 2 with the previous results [7,8]. In Fig. 4 the values of are normalized by the applied out-of-plane strain, e^. They compare well

with results of Wang and Crossman [7], who used a similar finite-element

model with constant strain triangles, and Pagano [8], who used an analytic

approach based on Reissner's variational principle. Our grid geometry was justified by this comparison. The approximation of the following elastic

properties were used to calculate

= 137.9 GPa (20 Msi)

E^ = 14.5 GPa (2.1 Msi)

Gg^^ = G^^ = 5.86 GPa (0.85 Msi) (1)

'’U

=

''tt

= 0-2'

The more correct elastic properties for glass-epoxy are listed in

Table 1, where and G^^ ^ G^^.. The influence of this approximation

on edge stresses is given in reference 9. Such approximations of elastic

properties are often made for a variety of elasticity solutions in composite

materials, and they can alter the model predictions significantly. In this

study we made no such approximation.

309
Stress distributions along the midplane for woven ([0/90]g^) and nonwoven ([0/90]^) laminates are compared in Fig. 5. The most significant

result is that is compressive near the edge (0.85 < y/B ^ 1.0) without the

thermal load. With the thermal load, the compressive stress, a^, increases.

For the nonwoven laminate, the opposite is true: is tensile near the edge

and increases in tension with the thermal load. Hence, the weave geometry

suppresses delaminations at the midplane.

The presence of a free edge also shows the symmetry of the stress

distribution. If no free edge existed at y/B = 1.0, the midplane distribution would be symmetric about y/B = 0.5. With the presence of a free edge

at y/B = 1 , the symmetry is skewed where the largest stress occurs at the

interior near y/B = 0. The largest tensile stress occurs in the interior

(0.15 < y/B < 0.8)

.

h ir Stress distributions along the midplane for [90/0]^ woven and ([90/0]^)

nonwoven laminates are compared in Fig. 6 . For the nonwoven laminate, is

compressive near the free edge, as expected. The stress near the edge in

the nonwoven laminate is in tension and decreases toward compression with the

addition of the thermal load AT = -100 K. Again, we observe a beneficial

effect of weave curvature on the edge stresses at low temperatures. The

symmetry of the distribution for [90/0]^ is skewed, decreasing near

the free edge and increasing in the interior near y/B = 0. The largest

compressive stresses occur at 0.4 < y/B < 0.6 for the [90/0]^ laminate.

The opposite is true for the [0/90]^^ laminate.

For both woven grid geometries, the weave geometry was beneficial at

low temperatures. The opposite was true for the [0/90]^ nonwoven laminate —

the thermal load increased the opportunity for delamination initiation at the

midplane.

Edge stress distributions through the thickness (0 ^ z/H S 1.0) for

woven ([0/90]g^, [90/0]^^) and nonwoven ([ 0 / 90 ]g, [90/013) laminates are compared in Figs. 7 and 8 . Again, we observe a beneficial effect of weave curvature on the edge stress at low temperatures for the [0/90]| laminate.

The opposite was true for the [ 0/9013 woven laminate where there was a large

tensile delamination stress, near z/H = 0.5, which increased with the

thermal load. The opposite was true at the midplane for the same laminate.

Hence, we expect delamination initiation at the 90/0 interface but not at the

midplane

.

All the above results assumed no weave curvature in the x-z plane.

Although this restriction is not a serious limitation for small weave angles,

it is most significant at the free edge for large angles. With curvature at

the edge in the x-z plane, we expect the stress in the [90/0lg to

decrease at z/H =0.5 and increase at z/H = 0 (the midplane). At the edge

of the [O/ 90 I 3 laminate, we expect to decrease at z/H =0.5 and z/H = 0.

To discover the exact influence of weave curvature in the x-z plane on edge

stresses requires a three-dimensional finite-element model. In this

study, we described only the effect of weave curvature in the y-z plane on edge stresses.

310
For the purpose of discussion, we selected extreme geometries to compare

general trends between woven and nonwoven geometries. Actual woven G-10CR

and G-11CR laminates have weave angles between 10° and 15°. Both [0/90]^^

and [90/0]g^ geometries exist at the edge of a real laminate. In Figs. 9

and 10, we compare edge stress distributions for these geometries. Again, we

see the beneficial effect of weave geometry at low temperatures, but expect

delaminations to occur at the midplane of the [0/90]^^ laminate. When the

warp angle, 6, increases to 45°, these midplane stresses become compressive.

The opposite trend occurs at the 0/90 interface of a [0/90]® laminate when 6 increases from 15° to 45°. Hence, the region of possible delamination

changes from the [0/90]® midplane to the [90/0]® interface as 0 increases.

O O

In all load cases discussed here, we assumed only a 100-K temperature

decrease for the thermal load, which is only one- third of the temperature

decrease experienced at an operating temperature of 76 K. Hence, the thermal

effects at 76 K clearly dominate the edge stresses.

Although the weave geometry is, in general, beneficial in reducing edge

stresses at low temperatures, we expect delaminations and transverse crack

formation in the fill-fiber bundles to occur at or near the laminate free

edge for higher thermal-mechanical loads. For some weave geometries it is

possible that damage may initiate in the interior. The heterogeneous nature

of the weave geometry and stress distributions at and near the free edge

explains why the damage at the edge is not representative of damage at the

interior

.

Conclusions

In general, the weave geometry is beneficial in reducing edge stresses

at low temperatures. For a [0/90]^ nonwoven laminate, the edge stress

normal to the midplane changes to compression with increasing warp angle.

For [0/90]® woven laminates, compressive edge stresses increase with decreasing temperature. For [90/0]| woven laminates, edge stresses at the midplane

are tensile and decrease to compression with decreasing temperature. For

both woven laminate types, [0/90]® and [90/0]|, the free edge skews the

symmetry of the midplane stress distribution by decreasing the magnitude of

the symmetric edge stress. This trend in skewed symmetry also occurs at low

temperatures. Unlike nonwoven laminates, edge stresses in woven laminates

are dominated by thermal loads. The larger warp angles demonstrated increased sensitivity of edge stresses to small changes in temperature. If

delaminations occur at the higher thermal-mechanical loads, initiation will

most likely occur in the interior. If delaminations occur at the edge,

initiation will most likely occur at the midplane of the [0/90]| laminates

and change to the 90/0 interface of the [90/0]| laminates with larger warp

angles. In general, smaller weave angles in the y-z plane are better at crack

suppression than laminates with larger weave angles. All conclusions are based on predictions where weave curvature is confined to the y-z plane.

Improved results are possible if the weave curvature is included in the x- z plane.

311
Acknowledgment

This study was sponsored by the Department of Energy, Office of Fusion

Energy.

References

1. M. B. Kasen, G. R. MacDonald, D. H, Beekman, Jr., and R. E. Schramm,

Mechanical, electrical, and thermal characterization of G-10CR and

G-1 1 GR glass/epoxy laminates between room temperature and 4 K, Advances

in Cryogenic Engineering , Vol. 26, Plenum Press, New York, pp. 235-244

( 1980 )

.

2. M. B. Kasen, Mechanical and thermal properties of filamentaryreinforced, structural composites at cryogenic temperatures, 1

Glass-reinforced composites. Cryogenics , Vol. 15, pp. 327-349 (1975).

3. M. B. Kasen, Cryogenic properties of f ilamentary-re inforced composites:

an update. Cryogenics , Vol. 21, pp. 323-340 (1981).

4. R. D. Kriz, Influence of damage on mechanical properties of woven

composites at low temperatures, J. Compos. Technol. Res. , Vol. 7, No. 2,

pp. 55-58 (1985).

5. R. D. Kriz and W. J. Muster, Mechanical-damage effects in woven laminates at low temperatures. Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures--VIII , NBSIR 85-3025, National

Bureau of Standards, Boulder, Colorado, pp. 49-86 (1985).

6. R. Byron Pipes and N. J. Pagano, Interlaminar stresses in composite

laminates under uniform axial extension, J. Comp. Mater. , Vol. 4,

pp. 538-548 (1970).

7. A. S. D. Wang and F. W. Crossman, Some new results on edge effect in

symmetric composite laminates, J. Comp. Mater. , Vol. 11, pp. 92-106

(1977).

8. N. J. Pagano, Stress fields in composite laminates, Int. J. Solids

Struct . , Vol. 14, pp. 385-400 (1978).

9. R. D. Kriz, Effect of material properties on interlaminar stresses in

angle-ply composite laminates. Report No. VPI-77-16, Virginia Polytechnic Institute and State University, Blacksburg, Virginia (1977).

312
Table 1 . Elastic-thermal properties [5]

PROPERTY

ISOTROPIC CONSTITUENTS TRANSVERSELY ISOTROPIC

COMPOSITE

BISPHENOL E-GLASS

EPOXY FIBER GLASS-EPOXY (V^ = 0.50)

Young’s modulus,

GPa

E = 7.16 E = 82.22

tu

ra

II

II .69

17.88

Shear modulus,

GPa

G = 2.77 G = 32.55 ^Slt = ^tt

= 6.801

6.888

Poisson's ratio V = 0.293 V = 0.263 II

II

-p 0. 278

0.298

Coefficient of — — = 5.963

thermal expansion, = 23*39

= volume fraction of glass fibers.

Subscript I indicates fiber direction; subscript t indicates transverse plane.

313
Figure 1. Geometries defined for woven and nonwoven symmetric laminates

loaded with uniform axial strain.

314
Nonwoven

Figure 2. Finite-element grids of woven and nonwoven [0/90]® laminates (0 = 0).

315
Figure 3- Complete finite element grids

woven laminates with complete

of [90/0]| and [0/90]®

symmetry showing both edges316
/*x.

MPa

160

0 0.2 0.4 0.6 0.8

y/B

CO

X

10 ^

N

1.0

Figure 4. Comparison of present finite element results with previous results

for 0 ^ stress along the midplane (z = 0).

317
15

y/B

Figure 5. Comparison of midplane stress distribution for [0/90]^^ woven and

[0/90]„ nonwoven laminates with and without thermal loads.

318
15

Figure 6

80- 1 1 P ^ jL 0°

H

90° ,

€x= 0.01

B =1.33 H

edge: [0/90]^

AT=-100 K

45®

CO Q.

N

-40 -

-80

CO

. N

--10

-15

Comparison of midplane stress distributions for [0/90]^^ woven and

[ 90 / 0 ]g nonwoven laminates with and without thermal loads.

319
Z/H

Figure 7.

c^z. ksi

-15 -10 -5 0 5 10 15

thermal loads.

320
ksi

-20 -15 -10 -5 0 5 10 15 20

Figure 8. Comparison of edge stress distributions through the thickness

for [90/0]g woven and [90/0]^ nonwoven laminates with and

without thermal loads.

321
,

MPa

y/B

Figure 9.

1 R Comparison of midplane stress distributions for [0/90]^

[90/0]g^ woven laminates with and without thermal loads.

and

322

,

ksi
ksi

-15 -10 -5 0 5 10 15

Figure 10. Comparison of edge stress distributions through the thickness for

[0/90]]^^ and [90/0]^^ woven laminates with and without thermal

loads.

323

TBCHNOLOOy TRANSFER

325

TECHNOLOGy TRANSFER PROGRAM

LEADER: N. J. Simon, NBS

CONTRIBUTORS: G. Latsis, P. Pollas, R. P. Reed

OBJECTIVES

• Organization of workshops to promote interaction between designers and

material specialists, to discuss issues related to low-temperature material

needs of the fusion energy program, and to present new low-temperature data

for structural alloys, composites, and weldments

• Preparation of an annual report as well as monthly highlight reports to

OFE/DOE.

• Evaluation of low-temperature mechanical and physical properties of

structural materials for superconducting magnets; preparation of handbook

pages and supporting documentation; distribution of handbook pages to

workshop attendees and others in the low-temperature materials applications

community; inclusion of pages in the "Materials Handbook for Fusion Energy

Systems"

RESEARCH PAST YEAR (1985)

• Preparation of "Materials Studies for Magnetic Fusion Energy Applications —VIII," (NBSIR 85- 3025 , 3^0 pages. May 1985, editor R. P. Reed)

• Preparation of handbook pages and supporting documentation covering

mechanical properties of AISI 304 (N, LN, HN, LHN); evaluation of physical

property data of AISI 304 alloys) and preparation of corresponding handbook

pages; submission of handbook pages for review to the "Materials Handbook

for Fusion Energy Sytems"

• Collection of data for handbook pages on CDA 102 copper and copper alloys

RESEARCH THIS YEAR (1986)

• Continued preparation for the low-temperature materials workshop to be

held October 6-8 1986.

• Preparation of "Materials Studies for Magnetic Fusion Energy Applications —IX," (NBSIR 86- 3050 , 338 pages, 1986, editor R. P. Reed)

• Establishment of an evaluated data base on CDA 102 copper and selected

copper alloy properties to be presented in the form of handbook pages,

evaluation of the data proceeds.

327

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iV

CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

Lawrence Berkeley Laboratory

University of California

Berkeley, California

October 2-4, 1985

This annual workshop provides an opportunity for discussion of problems,

advances, and goals in the development of low temperature materials for

superconducting magnets for fusion energy systems. The theme of this year’s

workshop was "Low Temperature Materials - Past Accomplishments and Future

Challenges." Emphasis was on the following:

• Recent accomplishments and lessons learned from the construction of the

Mirror Fusion Test Facility (MFTF-B) and the Large Coil Task.

• Identification of materials needs for future magnet system.

• New low temperature materials developments in the United States and Japan

and international research cooperation.

329
PROGRAM

Cryogenic Structural Materials Workshop

Lawrence Berkeley Laboratory, University of California

Berkeley, California

October 2—4, 1985

WEDNESDAY, OCTOBER 2, 1985

OVERVIEW AND RECENT PAST EXPERIENCES

9:00 a,m.

9:20 a,m,

9:50 a.m,

10:20 a.m. 11:10 a.m.

Overview of Cryogenic Materials Needs

Lessons Learned in Construction of a Large Number of Superconducting Magnets

for MFTF-B

1. A Laboratory Viewpoint

2. A Fabricator’s Viewpoint

3. Private Industry’s Viewpoint

Future Needs for Cryogenic Structural

Materials Development — Tokamaks and

Tandem Mirrors

V. Der, OFE

J. Zbasnik, LLNL

R. Bonina, CBI

R. Baldi, GD/C

E. N. C. Dalder,

LLNL

STRUCTURAL ALLOY EVALUATION AND DEVELOPMENT ACTIVITIES

1:30 p.m.

2:00 p.m.

2:30 p.m.

3:00 p.m.

3:45 p.m.

4:15 p.m.. 4:30 p.m.

Status of Evaluation of Cryogenic Structural

Alloys in the NBS Program

Evaluation of Candidate ICCS Conductor

Sheath Alloys

Modification and Development of Candidate

ICCS Conductor Sheath Alloys

Cryogenic Properties of Advanced Aluminum

Alloys

Relationships Between Processing, Microstructure, and Properties of Binary Fe-Mn

Austenitic Alloys

Joint Collaborations in Development and

Evaluation of New Cryogenic Structural

Materials

1. The Japanese Prospective

2. The U.S. Prospective

R. P. Reed, NBS

L. Summers, LLNL

R. Ballinger, MIT

J. Glazer, LBL

M. Strum, LBL

J. W. Morris, LBL

R. P. Reed, NBS

330
THURSDAY, OCTOBER 3, 1985

FABRICATION TECHNOLOGY FOR CRYOGENIC STRUCTURAL ALLOYS

9:00 a.m. Overview of Welding Development Needs for E. N. C, Dalder,

Future Cryogenic Magnet Case Materials LLNL

9:30 a .m. Development of Filler Metals for Next K. Banks, Teledyne

Generation of Cryogenic Magnet Case McKay

Materials T. A. Siewert, NBS

10:45 a »m

,

Fusion Welding of Fe-Mn Austenite Alloys J. Chan, LBL

11:15 a.m. Cryogenic Mechanical Properties of Welded M. Shimada, Kobe

22Mn-13Cr-5Ni Steel Steel

11:45 a.m. Mechanical Properties of High Manganese s. Onedera and

Austenitic Steels and Their Welds at R. Miura, Japan

Cryogenic Temperatures Steel Works

NONMETALLIC STRUCTURAL COMPOSITES FOR CRYOGENIC SERVICE

2:00 p . m

.

The Need for High Elastic Modulus, Low CTE E. N, C, Dalder,

Composites in Future TMR Choke Coils LLNL

2:30 p.m. Progress in Cooperative Irradiation Testing

of Composites

M. B, Kasen, NBS

3:00 p . m

.

Progress in Model Development and

Verification for Composites

R. D. Kriz, NBS

3:30 p.m. Interaction of Face Pressure and Interlaminar Shear on Thin Organic Laminates

H. Becker, MIT

4:00 p.m. Use of the NLTNIF for Magnet Materials

Research

R. Coltman, ORNL

FRIDAY, OCTOBER 4, 1985

FATIGUE, CREEP, AND PROCESSING EFFECTS ON PROPERTIES

9:00 a.m. Threshold Effects in Cryogenic Fatigue

Crack Growth

R. L. Tobler, NBS

9:30 a.m. Creep in Cryogenic Structural Materials R. P. Reed, NBS

10:00 a.m. Time Dependent Deformation in Austenitic

Stainless Steels at 4 K

L. Roth, Wes t inghousc

10:45 a.m. Processing Effects on Tensile and Fracture E. N. C. Dalder,

Toughness of Type 304LN Stainless Steel

at 4 K

LLNL

11:15 a.m. Processing Effects on Tensile and Fracture

Toughness of Type 316LN Stainless Steel

at 4 K

P. T. Purtscher, Nl'

331
FRIDAY, OCTOBER 4, 1985, continued

1:30 p.m.

1:30 p.m.

3:00 p.m.

3:00 p.m.

CONDUCTOR AND COIL TASK GROUP MEETINGS

Identification of Heat Treatment Facilities

for ICCS Sheath Materials

Identification of Heat Treatment "Windows"

for Alloyed Nb^Sn Conductors

Identification of Structural Failure Modes

for ICCS Conductors and Coils

Identification of Heat Treatment "Windows"

for ICCS Sheath Alloys

M. 0. Hoenig, MIT,

Chair

D, C. Larbalestier

,

University of

Wisconsin, Chair

V, Der, OFE, Chair

E, N, C. Dalder,

LLNL, Chair

332
CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

SUMMARY OF TECHNICAL PRESENTATIONS

Session 1 - Overview and Recent Past Experiences

Projected cryogenic materials needs were outlined by V. Der (OFE)

.

The importance of data on high performance copper alloys for the Compact

Ignition Tokamak (CIT) was stressed. This device will operate with liquid

nitrogen cooling, so properties between 77 K and room temperature need to be

determined. Other material data needs for superconducting magnet development

included information on the ICCS sheath alloys (e.g., 900-series Inconels),

316 LN, and weldments. The need for data on irradiation effects, creep, and

standardized methods of testing was also discussed. The importance of balance

between near-term and long-term DOE fusion energy program goals in setting

priorities for materials studies was stressed.

A summary of lessons learned from the fabrication of superconducting

magnets for MFTF-B was also presented in this session. A detailed account was given by R. Bonina (CBI) of how welding procedures were developed to ensure

high quality welds of thick sections in various orientations. Problems of

alignment and clearance between the large-scale individual components that make

up the entire tandem mirror machine of MFTF-B were some of the unusual aspects

of construction discussed by R. Baldi (GD/C).

E. N. C. Dalder (LLNL) spoke on future material needs for both tokamaks

and tandem mirrors. The large forces expected in current tokamak designs were

stressed; examples are a centering force (directed toward the center post of

the tokamak) of over 300 MN for INTOR and stresses of about 700 MPa for the

magnet case for TIBER. The need for data on irradiation effects in insulators,

creep of structural alloys, and the effects of the large number of cycles to

which these machines will be subjected were also discussed.

Session 2 - Structural Alloy Evaluation and Development Activities

R. P. Reed (NBS) discussed factors influencing cryogenic fracture

toughness of austenitic steels, particularly where Mn replaces Ni, and also

where Mo and Interstitial C+N additions contribute to increasing the yield

strength. The evidence for inclusion effects and the expected magnitude of

these effects were examined. New studies on effects of C, N, Ni, Mo, and Mn

content on AISI 317LN were outlined. L. Summers (LLNL) presented an evaluation of Inconels 903, 905, and 909 as candidate alloys for the ICCS conductor

sheath. These alloys must withstand extensive heat treatment because the Nb^Sn

superconductor has to be fabricated situ , owing to its brittleness. Inconel

903 has adequate il-K toughness (150 MPav^m) and yield strength (750 MPa) but a low Young's modulus after heat treatment. The potential cryogenic use of a new aluminum alloy with Li additions (2090) was discussed by J. Glazer (LBL). Both

toughness and yield strength for this alloy increase with decreasing temperature in contrast to other aluminum alloys.

333
This session concluded with a discussion by J. W. Morris (LBL) and

R. P. Reed of collaboration between the United States and Japan on the

development and testing of new cryogenic alloys, such as high-Mn, high-N

austenitic steels and high-purity, Mo-strengthened austenitic steels.

Collaboration on test standards and structural design standards is also

occurring. Details of the exchange of materials and the test program in the

United States at NBS, LLNL, and LBL were explained.

Session 3 ~ Fabrication Technology for Cryogenic Structural Alloys

In an overview of welding development needs for cryogenic magnet case materials, E. N. C. Dalder noted that two types of welding processes, each with

^-K qualified consumables, were desirable. A welding process with a high

deposition rate (SA, FCA, or ES) that could be automated to maximize productivity would be appropriate for many situations. But operator-dependent, highintegrity welding processes (e.g., GMA-P, SMA, or GMA) would be necessary for

vertical, overhead, and other difficult welds. Nondestructive inspection

methods must be validated before construction: for example, it was impossible

to inspect thick section (7 - 13 cm) welds on the yin-yang magnets for MFTF-B.

The U.S. goals for 4-K weld properties formulated during the 19811 Gryogenic

Materials Workshop in Vail, Colorado, were reiterated:

4-K Yield Strength Il-K Fracture Toughness

Automated process 1000 MPa

Operator-dependent process 1000 MPa

150 MPa/m

200 MPa/m

Details of the LLNL and LBL welding program with 31 6L and other consumables

were also presented.

T. A. Siewert (NBS) discussed his analysis of the cryogenic weld data

base. In particular, he showed that Charpy V-notch impact energy at 77 K

cannot at present be well correlated with i1-K fracture toughness, though the

77~K lateral expansion measured in the Charpy test is a somewhat better

predictor of ii-K yield strength and toughness. J. Chan (LBL) gave results of

weldment properties produced by gas shielded welding (75/S He, 255S Ar) of Fe-Mn

austenitic alloys. The welding program goals for yield strength were met, but

low fracture toughness values had not been overcome yet. M. Shimada of Kobe

Steel discussed electron-beam welding (EBW) of a promising alloy under

development: Fe-22Mn-1 3Cr-5Ni-0. 22N (0. 01 3P » O.OQilS, 0.3^131). Toughness values

of above 250 MPa/m at 11 K fall within the JAERI goals for Japanese fusion

systems, and no porosity or hot cracking of the weldment was observed. Some

decrease in N content (about 0.017 wt.^) occurred, which decreased the yield

strength. About twenty tests have been made. The effect of different

welding processes (SMA, FCA, GTA) on toughness of high-Mn austenitic steels was discussed by R. Miura (Japan Steel Works). Charpy impact energy data at 77 K were presented.

Session M - Nonmetallic Structural Composites for Cryogenic Service

The need for high-modulus, low- thermal-expans ion coefficient composites in

future tandem mirror choke coils was described by E. N. C. Dalder. Such coils

334
may have fields ranging from 25 to 50 T, and a new superconductor, PbMo^Sg

(H^ ~ 60 T) may be used. This strain-sensitive superconductor will require a high-modulus nonmetallic such as a graphite filament aluminum matrix composite

or the sputtering of the high-H^ superconductor onto a spiral grooved reinforcement plate with Kapton deposited as insulation.

M. B. Kasen (NBS) discussed procedures for fabricating small, wellcharacter ized specimens to be used in screening nonmetallics for low temperature radiation resistance. The procedure involves vacuum degassing and

impregnation. Resins under consideration in a cooperative testing program

include the standard G-11CR epoxy system and bisphenol polyimide (a state-ofthe-art resin). A type-E-glass fiber reinforcement with a proprietary finish

is being tested. R. D. Kriz (NBS) described modeling studies based upon a unit

cell of a woven reinforced composite. Results of compression and torsion tests

on a polyimide woven glass composite of 1 -mm thickness were also presented at

this session. These studies were carried out by H. Becker (MIT) who was unable

to attend the workshop.

The use of the National Low Temperature Neutron Irradiation Facility

(NLTNIF) at ORNL for magnet material research was the subject of a presentation

by R. Coltman (ORNL). This facility will enable both irradiation and testing

to be carried out at 4 K without warmup, but specimen size limits and irradiation heating rates that affect the internal specimen temperatures need to be

considered in experiment design. The facility, to be operational in 1986, is

the first of its kind.

Session 5 - Fatigue, Creep, and Processing Effects on Properties

Because tokamak fusion energy machines will be operated in pulsed modes,

threshold fatigue crack growth rate (FCGR) data is an important consideration

in efficient structural design of the superconducting magnets. R. L. Tobler

(NBS) described an automated apparatus that has been used to obtain FCGR data

on AISI 300 series austenitic steels at 4 K that are several orders of magnitude below rates of previously published measurements. There was considerable

discussion about the significance of the preliminary results for design

purposes

.

Recent results have indicated that creep may occur at cryogenic temperatures in copper and austenitic steels. R. P. Reed reported results on AISI 316

and 310 at 76 K that showed that no steady-state creep occurred and that

probably a minimum stress would be necessary to initiate creep. Because of

the complexity of designs for structural members of magnets, stresses can approach yield strengths in some parts of the structure, so such data are important. Further discussion on possible athermal mechanisms for cryogenic

creep, such as quantum mechanical tunneling was given in the next talk by

L. Roth ( Westinghouse)

.

A presentation by P. T. Purtscher (NBS) on processing effects on 4-K t*-nsile and fracture toughness concluded that grain size affected yield strength

and fracture toughness inversely, as does nitrogen content, and that rearrcinging inclusions, precipitates, or their distribution in the microstruoture

through hot-working could improve toughness.

335
Session 6 - Conductor and Coil Task Group Meetings

The main topic of discussion was the extent of allowable heat treatment

for conductor sheath alloys, such as Inconel 905 and 903, that would permit in

situ processing of the Nb^Sn superconductor but not degrade the iJ-K mechanical

properties of the alloy.

A discussion of failure modes and design philosophy for superconducting

magnets under static and cyclic loading concluded this session.

336
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‘
INFORMATION RETRIEVAL SERVICES

Numerous requests for material property data are received by our division each year. During the past year, we completed data searches and supplied the following material properties to other laboratories and companies:

Information Provided to

Kfc for Nitronic 33 at 4 K national laboratory

measurement techniques at 4 K institute of technology - Japan

measurement techniques at 4 K atomic reactor company

magnetic susceptibility of austenitic

steels at 4 K

research institute - Japan

tensile properties of Au from 4 to 10 K

for a space application

aerospace company

axial fatigue and notch tensile strength

of A1 6061 T6 and A1 7075 T73 at 20 K

for an engine to be used in space

aerospace company

physical and mechanical properties of

Ti and A1 alloys at 3 K for an electrooptical space application

instrument company

thermal conductivity of CDA 102 Cu and

A1 5657 from 4 K to room temperature

aerospace company

mechanical properties of 304 alloys at

cryogenic temperatures for cryogenic

tankage

cryogenic equipment manufacturer

mechanical properties of 21-6-9 weldments

at 4 K for a particle accelerator

national laboratory

tensile properties, Kj^, and magnetization

of high-N austenitic stainless steels

university research center

properties of nonmetallics university research center

load-controlled fatigue of CDA 101-104 Cu

at 76 K

electrical equipment manufacturer

3.5 Ni steels - mechanical properties of

base and welded material at cryogenic

temperatures

national laboratory

337
INFORMATION RETRIEVAL SERVICES, continued

Information Provided to

magnetic permeability and for Nitronic

33 at K

national laboratory

cryogenic mechanical and thermal properties

of Pt for a temperature monitor

(space application)

manufacturer of aerospace

components

cryogenic properties of 304 and 316 alloys

and Cu

national laboratory

cryogenic properties of polymers and

structural alloys

university research center

cryogenic properties of structural alloys materials laboratory - France

PVT properties of O

2 and N2 for highpressure pneumatic systems (military

radar applications)

cryogenic equipment manufacturer

of 304 and Ti-6A1-4V alloys at 4 K

for a particle accelerator

aerospace company

cryogenic properties of 304, 316, and

21-6-9

aerospace company

fracture properties of A1 2219 T62 at 20 K aerospace company

thermal expansion of Fe-l8Cr-l8Mn to 4 K university research center

1) tensile strength, yield strength, shear

strength, elastic modulus, and fracture

toughness of eleven structural alloys,

G^IOCR, and G-11CR from 300 to 4 K; 2) thermal expansion, thermal conductivity,

and specific heat data for structural

alloys from 300 to 4 K; (cryostat design for particle accelerator)

national laboratory

338
u.». OtPT. OF COMM. 1. PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No. 3. Publication Date BIBLIOGRAPHIC DATA

SHEET (See instructions) NBSIR 86-3050 May 1986

4. title and subtitle

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures - IX

5. authoR(S)

R, P, Reed, Editor

(. PERFORMING organization (I[ joint or other than NBS. see instructions;

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

7. Contract/Grant No.

S. Type of Report & Period Covered

9, SPONSORING organization NAME AND COMPLETE ADDRESS (Street. City. State, ZIP)

Department of Energy

Office of Fusion Energy

Washington, DC 20545

10.

SUPPLEMENTARY NOTES

Previous reports in this series: NBSIR 85-3025, NBSIR 84-3000, NBSIR 83-1690,

NBSIR 82-1667, NBSIR 81-1645, NBSIR 80-1627,

NBSIR 79-1609, NBSIR 78-884

Document describes a computer program; SF-185, FlPS 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 results of a research program to produce material property

data that will facilitate design and development of cryogenic structures for the

superconducting magnets of magnetic fusion energy power plants and prototypes. The

program was developed jointly by the staffs of the National Bureau of Standards and

the Office of Fusion Energy of the Department of Energy; it is managed by NBS and

sponsored by DOE. Research is conducted at NBS and at various other laboratories

through subcontracts with NBS.

Research results for 1985 are summarized in an initial "Highlights of Results"

section and reported in detail in the technical papers that form the main body of

this report. The technical papers are presented under four headings reflecting the

main program areas: Welding, Nonmetallics , Structural Alloys, and Technology Transfer

Objectives, approaches, and achievements are summarized in an introduction to each

program area.

12.

KEY WORDS (Six to twelve entries; alphabetical order; capital! ze only proper names; and separate key words by semicolon S/ austenitic steels; composites; cryogenic properties; elastic properties; low temperatures; mechanical properties; nonmetallics; stainless steels; structural alloys; welds

13.

AVAILABILITY

Unlimited

1 1 For Official Distribution. Do Not Release to NTIS

Order From Superintendent of Documents, U.S. Government Printing Office. Washington. D.C.

20402.

[25 Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

348

IS. Price

jS'OMfc-i.CC *

'

U. S, GOVERNMENT PRINTING OFFICE 1986 67? ? 'n a* m

I