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|Siii°iF STANDARDS & TECH R.I.C. NBSIR 87-3067

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MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES -X

To

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

Prepared by

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

QC

100

- U 56

87-3067

1987
IBBDDDB9IIIni
NBSIR 87-3067 UC 20b, c

TECHNICAL REPORTS

Research Ijilomiation Center

A

T alionai Bureau of Standards

Gaithersburg, Maryland 20899

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY - U:

APPLICATIONS AT LOW TEMPERATURES -X ***+

Edited by

R.P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

May 1987

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
II

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CONTESTS

SUMMARY v ORGANIZATIONAL CONTACTS vi

PROGRAM DESCRIPTION vii

STRUCTURAL ALLOYS

Structural Alloys Program 3

Summary of Mechanical Property Measurements of Structural Alloys ... 5 Tensile Strain-Rate Effects in Liquid Helium 9

R. P. Reed and R. P. Walsh

Modification of 316LN-Type Alloys 31

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

Molybdenum Effect on Fe-Cr-Ni-Alloy Elastic Constants 43

H. M. Ledbetter and S. A. Kim

Molybdenum Effect on Volume in Fe-Cr-Ni Alloys 55

H. M. Ledbetter and M. W. Austin

Low-Temperature Manganese Contributions to the Elastic Constants

of Face-Centered-Cubic Fe-Cr-Ni Stainless Steel 67

H. M. Ledbetter and S. A. Kim

Strength and Toughness of AISI 304 and 316 at 4 K 77

N. J. Simon and R. P. Reed

Properties of CDA 104, 155, and 175 Copper Alloys . 83

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

Deformed-Polycrystalline-Copper Elastic Constants 127

H. M. Ledbetter and M. W. Austin

Low-Temperature Elastic Constants of Deformed

Polycrystalline Copper 139

H. M. Ledbetter and S. A. Kim

Acoustoelastic Residual-Stress Measurements:

Role of Anisotropic Dislocation Arrays 151

H. M. Ledbetter

Temperature Variation of the Elastic Constants

of Aluminum Alloy 2090-T81 163

J. Glazer, J. W. Morris, Jr., S. A. Kim,

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

iii
WELDING

Welding Program ...... 169

Summary of Mechanical Property Measurements of Welds ........ 170

The Role of Inclusions in the Fracture of Austenitic

Stainless Steel Welds at 4 K ................... . 173

T. A. Siewert and C. N. McGowan

Cryogenic Strength, Toughness, and Ferrite Content

of Stainless Steel Welds ...... 185

C. N. McGowan

Influence of Molybdenum on the Strength and Toughness

of Austenitic Stainless Steels for Cryogenic Service ........ 285

C. N. McCowan, T. A. Siewert, and E. Kivineva

High-Energy-Beam Welding of Type 316LN Stainless Steel

for Cryogenic Applications . .......... 299

T. A. Siewert, D. Gorni, and G. Kohn

NONMETALLICS

Nonmetallics Program . . ........ 319

Laminated Plate Analysis for Low-Temperature Applications:

Copper Reinforced with Metallic and Nonmetallic Laminates 321

R. D. Kriz

High Quality Organic Matrix Composite Specimens

for Research Purposes ............ 343

M. B. Kasen

Strategy for the Data Base Construction on RadiationResistant Cryogenic Composite Insulators for Magnetic

Fusion Energy Applications .......... 347

M. B. Kasen

TECHNOLOGY TRANSFER

Technology Transfer Program 359

Ninth Annual Cryogenic Structural Materials Workshop 361

UNITED STATES-JAPAN COOPERATIVE PROGRAM

United States-Japan Cooperative Program 373

United States-Japan Low-Temperature Structural Materials

and Standards Workshop 375

Proposed Standard Method for Tensile Testing of

Structural Alloys at Liquid Helium Temperature . 387

Proposed Standard Method for Fracture Toughness Testing

of Structural Alloys at Liquid Helium Temperature . 405

Fracture, Acoustic Emission, and Adiabatic Heating of Austenitic

Stainless Steels at Liquid Helium Temperature 423

R. L. Tobler, T. Shoji, H. Takahashi, and K. Ohnishi
SUMMARY

This report contains the results of a research program

to determine the properties of materials that may be used in

crvogenic structures for the superconducting magnets of magnetic

fusion energy power plants and prototypes. Its purpose is to

facilitate their design and development. The program was developed jointly by the staffs of the National Bureau of Standards

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

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

at NBS and at other laboratories through subcontracts with NBS.

Research results for 1986 are presented in technical papers

under five headings that reflect the main program areas: Structural Alloys, Welding, Nonmetallics , Technology Transfer, and

I'nited States-Japan Cooperative Program. Objectives and research

highlights are summarized in the introduction to each program

area

.

Note : Certain commercial equipment, instruments, or materials

are identified in this report to specify the experimental procedure adequately. In a few instances, company names are used

to identify the source of specific research. In no case does

such identification imply recommendation or endorsement by the

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

or equipment identified is necessarily the best available for the

purpose

.

v
ORGANIZATIONAL CONTACTS

Specific technical questions may be directed to the following

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

fiscal year 1987.

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

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

National Bureau of Standards, Boulder, Colorado» 80303

Program Manager R. P. Reed (303)

(FTS)

497-3870

320-3870

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

H. 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

Welding C. N. McCowan (303) 497-3480

T. A. Siewert (303) 497-3523

Nonmetallics M. B. Kasen (303) 497-3558

R. D. Kriz (303) 497-3547

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

N. J. Simon (303) 497-3687

vi
PROGRAM DESCRIPTION

The overall objective of the program is to assist in the design,

construction, 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

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

MFTF-B, we are aware of new problems as they arise. This year we have

included for the first time results of research on copper and copper

alloys, in support of Compact Ignition Tokamak designs by Princeton and

Massachusetts Institute of Technology. Research projects contributing

to the second objective are described in the technical papers. Again,

research toward the establishment of test standards was emphasized.

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

by the series of workshops on Materials at Low Temperatures. Since

1982, handbook pages presenting the available data for specific materials have been distributed to members of the magnetic fusion energy

community who are involved with low- temperature materials.

vii

STRUCTURAL ALLOYS

1

STRUCTURAL ALLOYS PROGRAM

LEADER: R. P. Reed

CONTRIBUTORS: M. W. Austin, S. A. Kim, H. M. Ledbetter, P. T. Purtscher,

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

OBJECTIVES

• Dove lopment of strong, tough structural alloys for use in superconducting

magnets

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

pulsed superconducting magnets

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

measurements and structural design

RESEARCH HIGHLIGHTS

• The effects of strain rate on the tensile properties of three austenitic

stainless steels at 4 K were examined. In the range 4.4 x 10' 6 s' 1 to

8.8 x 10' 3 s' 1

, strain rates less than 2.2 x 10" 3 s" 1 had no effect on

tensile properties, and strain rates of 4.4 x 10" 3 s" 1 or larger reduced

the ultimate tensile strength. Estimates of the work put into the specimen during deformation, the stored energy in terms of dislocations and

dislocation interactions, and the dissipated heat were calculated. The

reduction of tensile strength was associated with specimen warming (to

100 K) , which was caused by the transition from nucleate to film-boiling

heat transfer on the specimen surface.

• Mechanical properties and electrical resistivity of three copper alloys

were measured at 4, 76, and 295 K. Effects of specimen orientation,

load- and displacement-control tensile testing, degree of hardness, and

grain structure and size were studied. Yield strengths were increased by

cold-rolling, but at the expense of electrical conductivity, ductility,

and toughness. Of the three alloys, CDA 104, 155, and 175, alloy 104 in

the hard condition had the best combination of yield strength and electrical conductivity.

• The elastic constants of the CDA 104 and 155 copper alloys were studied

semicontinuously between 295 and 4 K by ultrasonic methods. At 295 K,

cold-rolling reduced the elastic stiffness of the alloys, reflecting

changes in texture, nonpinned dislocation density, and anisotropic

dislocation array; after this plastic deformation, the elastic constants

showed time -dependent recovery. The dislocation-array anisotropy

destroyed the usual Cjjkl =

^ijlk macroscopic symmetry for the elastic

shear stiffness, suggesting large, fictitious internal strains. Between

3
295 and 4 K, the elastic-constant temperature dependence of the deformed

coppers approximated that of annealed texture -free copper.

• The effects of molybdenum and manganese on the elastic constants and

volume of Fe—Cr—Ni alloys were studied. Molybdenum alloying caused

unusual increases in volume that were attributed to changes in interatomic bonding. Molybdenum decreased the shear and Young's moduli and

increased the bulk modulus and Poisson's ratio. Manganese produced

predictable effects: it increased volume and decreased all the elasticstiffness constants. For manganese, the 4-K alloying effect resembled

the 295-K effect.

4
> MM ARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS

Country Fracture Fatigue Crack Fatigue * Alloy of Tensile Elastic Toughness Growth Rate S-N Report Volume

( Da 1 gnat ion

Supplier 295-4 K 295-4 K 295, 76, 4 K 295, 76, 4 K 4 K (page)

AUSTENITIC STEELS

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

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

Fe-1 3Cr-19Mn USSR III VI III (79) VI (41)

Fa-l 3Cr-22Mn-0 . 2N Japan VI VII (65)

Fa-l6Cr-8Ni-8Mn U.S. VIII VIII VIII (151,181)

Fa-l7Cr-18Mn-0.5N U.S. VII, VIII VII (65) VIII VIII (181)

F*-l7Cr-9Ni-8Mn U.S. III III (91)

Fe-l7Cr-l 3Nl-2Mo U.S. I, II, VI I, II, II, VIII, I I (15,71)

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

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

IX IV (147) V (185) VI (157) VIII (181,209,251)

IX (15,27,43)

Fe-17Cr-l 3Nl-2Mo U.S. II, IX II, IX II II (79)

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

Fe-18Cr-3Ni-l3Mn U.S. I, VI, III I, VIII I, VI, VII I (93) VIII III (91) VI (53) VII (85) VIII (167,219)

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

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

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

(304] III, IV, III, IV, III, IV, II (79,149,175)

VI, VII, VI, VII IX III (15,91,105,117

IX IV (37,101,203,

215,227)

V (71)

VI (73) VII (13,47,75,157)

IX (27)

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

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

IX IV (37,101,131

215,302)

VIII (181,229,251)

IX (27,141)

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

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

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX).

5
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

Alloy

[Designation]

Country

of

Supplier

Tensile

295-4 K Elastic

295-4 K

Fracture

Toughness

295, 76, 4 K

Fatigue Crack

Growth Rate

295, 76, 4 R

Fatigue

S-N

4 K Report Volume(page)

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

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

VIII (181)

IX (53)

Fe-19Cr-9Ni-N U.S. 1,111, IV I, III, IV, I, III, IV I (93)

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

215)

IX (27)

Fe-19Cr-9Ni-N U.S. II, III, IV, VIIII, III, IV, II, III, IV, II (35,79)

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

VIII, IX IV (37,101,203)

V (29) VI (113) VIII (181,207,229)

IX (159)

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

(0 . 03-0 . 29 )N- VIII (145,181)

(0.03-0.09)0 IX (159)

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

[216]

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

[304HN]

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

VI (113)

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

III (91) VI (157) VIII (151,181)

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

II (79) III (91)

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

[310] VII, V (71,145)

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

IX (95)

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

[200-300 grades]

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

VI V (171) VI (11,127,157)

VII (103)

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

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

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX).

6
MM A R Y • MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS, continued

A1 loy

( Designer Ion

Country

of

Supplier

Tensile

295-4 K Elastic

295-4 K

Fracture Toughness

295, 76, 4 K

Fatigue

Growth

295, 76,

Crack

Rate

, 4 K

Fatigue

S-N

4 K Report Volume

(page)

ALUMINUM ALLOYS

Al-6Cu-0.2Mn-0.lFe

(5083-0)

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

Al-0.6fe-0.2Cu0.131

(1100)

U.S. IX IX (107)

Al-AMg

(5083-0)

U.S. IV II IIIV

(35)

(185)

COPPER ALLOYS

99 . 99Cu-Ag

(CDA 101,102)

U.S. IX IX (95,107)

SUPERCONDUCTORS

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

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

KBSH 74-lbOq (II); NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-1690 (VI)'; NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX).

7
.'


TENSILE STRAIN-RATE EFFECTS IN LIQUID HELIUM*

R. P. Reed and R. P. Walsh

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The effects of strain rate on tensile properties of three

austenitic stainless steels at 4 K were examined. Strain rates

ranged from 4.4 x 10' 6 s' 1 to 8.8 x 10" 3 s' 1

. Strain rates less

than 2.2 x 10* 3 s' 1 had no effect on tensile properties. Strain

rates of 4.4 x 10' 3 s' 1 or larger reduced the ultimate tensile

strength, and stress—strain curves and temperature measurements

indicated specimen warming to 100 K. Calculations are presented

to estimate the work put into the specimen during deformation,

stored energy in terms of dislocations and dislocation interactions, and dissipated heat. The reduction of tensile strength

was associated with specimen warming, which was caused by the

transition from nucleate to film-boiling heat transfer on the

specimen surface.

INTRODUCTION -

Internal specimen heating effects during tensile tests at 4 K have been

questioned for a long time. The very low specific heat and thermal conductivity of metals at low temperatures lead one to anticipate near-adiabatic

conditions. Discontinuous yielding, a phenomenon in which very high strain

rates are experienced in local regions (similar to Luders band regions in

steel at ambient temperatures) , complicates the question, and the answer is

prerequisite to the establishment of test standards.

Read and Reed 1 studied discontinuous yielding during tensile tests of

alloy 304L at 4 K. Using embedded thermocouples and superconducting filaments in contact with the (flat) specimen surfaces, they showed that specimens heated locally to about 50 K during discontinuous yields. The yields

occurred in bands, about 10' 2 m wide, across the specimen. The specimen

extension corresponding to a typical load drop was about 1.3 x 10" 4 m. The

temperature rise, confined to the strained regions, was adequately described

by assuming adiabatic conditions, that is, all mechanical work was converted

to heat. The temperature returned to 4 K following each discontinuous

yield

^Submitted to Advances in Cryogenic Engineering Materials—vol. 34.

9
' Ogata and Ishikawa2 studied the effects of strain rate on the tensile

properties of austenitic stainless steels at 4 K. Strain rates were varied

from 8.3 x 10" 5 to 8.3 x 10" s' 1 on 6 . 25-mm-diameter round specimens. They

reported that the yield strength of alloy 304L was insensitive to strain

rate; at strain rates less than 3 x 10* 3 s' 1

, its ultimate strength was also

insensitive to strain rate; at strain rates above 3 x 10~ 2 s" 1

, the ultimate

strength decreased. They also measured the internal specimen temperature

during deformation by inserting a thermocouple into a small hole drilled

longitudinally into the specimen. Thermal spikes, associated with discontinuous yielding, occurred at the lower strain rates, and the magnitude of

the thermal spikes increased with increasing strain rate. At the highest

strain rate (3 x 10' 2 s' 1 ), the temperature rose continuously from 4 to

almost 150 K at failure. Temperature profiles for alloys 304L and 310 were

similar, but larger thermal spikes were observed in alloy 304L. Temperature

rises of about 50 K occurred during discontinuous yielding.

The same temperature rise during discontinuous yielding was observed

by Dobson and Johnson. 3 For alloy 304L, they reported a steady decline of

ultimate strength over their strain rate range, 3 x 10" 5 to 3 x 10” 3 s' 1

, and a linear decrease of yield strength when it was plotted versus the log

of strain rate.

Yamagami et al. 4 recently reported the tensile strain-rate sensitivity

at 4 K of three Fe—21Mn—14Cr—7 . 5Ni (wt.%) austenitic alloys with varying

N content (0.02 to 0.20 N) and grain size (42 to 284 /xm) . They measured

temperature rises with embedded thermocouples at strain rates from 1.2 x

10* 4 s' 1 to 1.2 x 10' 2 s' 1

. Sudden spikelike temperature increases (up to

about 40 K) were found at lower strain rates (< 1.2 x 10' 4 s' 1 ). At higher

strain rates, .the temperature at the thermocouple rose and fell much more gradually, but in larger increments. At the larger strains and the highest

strain rates, temperatures over 60 K were recorded, and the minimum temperature remained above 20 K. Associated with internal specimen heating,

the ultimate strengths decreased at the higher strain rates (8 x 10' 3

, 11.2 x 10' 2 s' 1 ). They found that tensile yield strength was insensitive

to low strain rates, but tended to decrease at strain rates greater than

5 x 10' 3 s' 1 if the yield strength exceeded about 1000 MPa. Two steels with

yield strengths of less than 900 MPa showed no change in yield strength at

higher strain rates. Thus, except for that of Dobson and Johnson, these

studies support the premise that the strain rate sensitivity of yield

strength depends on the magnitude of the yield strength: for stronger

alloys, yield strengths decrease at higher strain rates.

This paper presents the results and analysis of the dependence of the

tensile properties of three austenitic stainless steel alloys on strain rate

at 4 K. Alloys 310, 316LN, and 304L were selected because of their varying

yield strengths and austenite stabilities. Alloy 310, containing 25% Cr

and 21% Ni, does not transform partially to body-centered-cubic martensite

during plastic deformation at 4 K. 5 Alloy 304, containing only 18% Cr and

10% Ni with low N and C, has the least stable austenite and transforms

almost completely to body-centered-cubic martensite during deformation at

76 K; at 4 K the transformation is not as complete. 5

10
EXPERIMENTAL PROCEDURE

Specimens of alloys 310 and 304L were machined from 1 . 9-cm-diameter

bar stock; specimens of alloy 316LN, from 5.0-cm-thick plate. Alloy 304L

was tested in two conditions: annealed at 1066°C for 25 min (labeled

304L-A) , and annealed at 1050°C for 45 min (labeled 304L-B) . Microstructures of alloy 304L are shown in Fig. 1. Two groups of alloy 31 6LN were

annealed at 1150°C for 1 h; 316LN-A had a larger grain size and inclusion

density than 316LN-B (see Fig. 2). The 316LN-A specimens, taken from

near the plate surface of the 5.0-cm-thick plate, had a N content of

0.167 wt

% ; 316LN-B, from the interior, had a N content of 0.149 wt.%.

Alloy 310 was tested in the partially sensitized condition; its microstructure is shown in Fig. 3. The chemical composition, hardness, and

grain size of the alloys are listed in Table 1.

Round, tensile specimens (see Fig. 4) were machined to a 6.35-mm

diameter along a gage length of 41.9 mm, increasing with a radius of

9.52 mm to 12 . 7 -mm-diameter threaded ends. The entire specimen length

was 71.4 mm .

The cryostat equipment has been described in reference 6. Tests

at 4 K were conducted in boiling liquid He. New, high-elongation straingage extensome ters that have been developed for use at low temperatures

were used for these measurements. These extensometers , which are capable

of sensitivities of 10" 6

, have three flexing Al beams whose geometry

enables sensing of specimen strain to elongations of over 30% at low

temperatures. Four 350-Q strain gages were mounted to each beam and

wired in a Wheatstone bridge configuration. A 5-V dc power supply was

used to supply the excitation. For maximum sensitivity (10‘ 6

) at low

strain, a set of three-beam strain-gage extensometers was also used for

yield-strength determination. Elongation was measured over a specimen

gage length of 3.8 cm.

Crosshead rates were varied from 0.001 cm-min -1 to 2 cm’min' 1

. Since measurements at the lowest strain rate took over four hours to

reach the yield strength, these tests were not continued to specimen

failure. Reported strain rates refer to the rate of plastic deformation

within the reduced section of the specimens.

Type E (NiCr—Constantan) thermocouples were attached to some specimens for temperature measurements during straining. These fine (0.25 -mm)

wire thermocouples were first varnished, then taped, and finally clamped

to the midsection of the specimen. They recorded specimen temperature

throughout the tensile test in liquid He. By a similar technique, the

same thermocouples were occasionally used to measure the temperature of

the specimen grips. Magnetic readings of deformed specimens were obtained

from a bar-magnet, torsion-balance, commercial magnetometer.

11
EXPERIMENTAL RESULTS

The dependences of ultimate and tensile yield strength on strain rate

at 4 K are shown in Figs. 5 and 6. Notice the abrupt decrease of ultimate

strength at strain rates greater than 2.2 x 10" 3 s' 1

. Yield strength

remained almost independent of strain rate; with increasing strain rate,

it increased slightly for alloys 310 and 316LN and decreased slightly for

alloy 304L.

Reduction of area and tensile elongation are shown as a function of

strain rate in Fig. 7. These ductility parameters were nearly independent

of strain rate. Increased necking (higher reduction of area) appeared only

at the higher strain rates. Tensile elongation gradually decreased with

increasing strain rate. However, the considerable data scatter in both

parameters weakens the evidence for clear trends.

For all alloys, the yield strength, elongation, and reduction of area were insensitive to strain rate. The ultimate strength of all alloys

decreased at the same high strain rate and showed similar trends at lower

rates

The stress—strain curves at differing strain rates are shown in Figs. 8 through 10. At the lower two strain rates, discontinuous yielding is prominent. At the highest strain rate, discontinuous yielding is less evident,

the curves have an erratic appearance, and reduction of ultimate strength is

apparent. Clearly, the nature of the deformation process has changed. The

onset of discontinuous yielding is dependent on both alloy and strain rate.

The dependences of ultimate strength on strain rate are similar despite

differing modes of austenite fracture. In alloys 304L and 316LN, microvoids

initiate at inclusion sites in the specimen interior and coalesce to form

cracks (Figs. 11 and 12); usually a shear lip connects this interior region

to the specimen surface to form some sort of "cup-cone" fracture morphology.

Alloy 310 was partially sensitized (Fig. 13d), and as Fig. 13 (a and b)

illustrates, its primary fracture path was intergranular. In Fig. 13c,

notice the very fine microvoid sites within the grain boundary; presumably

these initiate at the carbides.

The temperature measurements of one thermocouple attached to the cen- tral portion of the specimen surface of alloys 316L and 304LN are shown

in Figs. 9 and 10. Notice the change from thermal spikes at lower strain

rates to thermal waves at the highest strain rate. The thermal spikes are associated with individual discontinuous yielding events. The maximum

temperature rise ranges from 50 to 85 K and varies with strain and strain

rate, in agreement with earlier work. 1 ’ 4 The thermal waves, which gradually

build up to temperatures as high as 96 K, are associated with the lack of

discrete discontinuous yields and with the wavy nature of the stress—strain

curves. The temperature rise of the surface, as measured here, is less than

the internal temperature measured by Ogata and Ishikawa: 2 at high strain

rates the internal temperature rises steadily to about 150 K; the surface

temperature variation is wavelike, and maximum temperatures approach about

100 K.

12
The transformed body-centered-cubic (a') martensite in austenitic

stainless steels is ferromagnetic; the parent austenite phase is paramagnetic. Magnetic readings were taken on deformed specimens to obtain relative amounts of a' martensite as a function of strain rate. The amount of

a' formation is depends on elongation. The normalized results (magnetic

readings per unit percent elongation) are reported as a function of strain

rate in Fig. 14. The results are surprising: in alloy 304L the relative

amount of a' martensite transformation increases with increasing strain

rate, but in alloy 316LN, a' decreases with increasing strain rate.

DISCUSSION

The work of deformation (E^ef) either remains stored in the specimen

(F. se ) or is dissipated as heat (Q) . Energy is stored primarily in dislocations created during plastic deformation and, in the process of work hardening, as interactive, dislocation stress fields. Heat is conducted from the

deformed volume, through the specimen (Qs )

and to the liquid He via the

specimen surface (Ql)

The purpose of these calculations is to describe specimen heating at

the higher strain rates, where ultimate strength suddenly decreases. In

another paper we will describe localized, discontinuous yielding. We want

to calculate or estimate the parameters E^ef. Egg, Qs , and Ql for tensile

tests at two strain rates (e) from the data of alloy 316LN (Figs. 5 and 9).

From the experimental results at the strain rate of 2.2 x 10" 3 s' 1

, we

assume that the average specimen temperature is 4 K. Thermal spikes are

associated with discontinuous yielding. But, because of the local nature

of the temperature spike, the specimen temperature relatively quickly re- turns to 4 K; the ultimate tensile strength is unaffected by these events.

At the strain rate of 4.4 x 10' 3 s" 1

, however, the stress—strain curves and

internal 2 and surface measurements indicate an overall increase of specimen

temperature, affecting the stress—strain relationship and ultimate tensile

strength

For these calculations it is useful to consider energy per unit time, or the derivative of E^ef. Egg, and Q as a function of time. From Ronnpagel

and Schwink, 7 E^ef = E^gf/e , Egg = Egg/e, and Q = Q/e . Then

Edef = ^se + Q> (!)

and energy is conserved at any instant of time.

Deformation Work

The work of deformation is the product of flow stress (<7f) , strain (e),

and volume of the deformed region (V)

The rate of work of deformation,

therefore, is

Edef = «"fV. (2)

The rate of work of deformation scales directly with flow strength, as seen

in Eq. 2. Higher strength alloys require more work. Using the strain rate

13
immediately preceding the detection of thermal waves in the deforming

316LN specimen, 2.2 x 10" 3 s' , and the ultimate tensile strength of

1650 MPa, we find that Ede f = 2.9 W. For the next highest strain rate

tested, 4.4 x 10' 3 s

1 (which produced thermal waves during deformation

and a reduced ultimate strength, 1520 MPa), Ede f = 5.5 W. For these

calculations, V is considered to be the entire specimen volume within

the gage length (2.54 cm).

Stored Energy

The stored energy may be estimated as the sum of the elastic strain

energy of the dislocations created during the deformation process and

the long-range elastic internal stresses. 7

Ese = NEdV + eafV/(2Y). (3)

Here, N is the rate of dislocation production, Ed

is the dislocation

strain energy, and Y is Young's modulus.

The strain energy of a dislocation is typically represented as:

Ed = (^b 2 /4?rk) In R/r0 = ;ub 2

, (4)

where /i is the shear modulus, b is the Burgers vector, k = 1 for screw

dislocations, and k = 1 —for edge dislocations (v = Poisson's ratio).

The term R/r0 is the ratio of the outer radius to the inner radius of

the cylinder (about the dislocation core) through which strain energy

is calculated.

The rate of energy to create dislocations, NEdV, is quite small.

The plastic deformation of the test to failure of 316LN specimens at a strain rate of 2.2 x 10' 3 s' 1 lasts about 200 s, and an estimate of the

maximum number of dislocations created during this test is 10 16 dislocations per m

2

. From the values V = 1.2 x 10' 6 m

3 for our tensile specimens, n = 81 GPa for alloy 316LN at 4 K, and b = 2 x 10' 10 m, NEdV is

about 0.2 W. The internal stress term is smaller, insignificant in these

calculations

From calorimetric measurements of deformed Cu at room temperature,

the stored energy represents from 5 to 10% of the deformation energy at

large strains. 8 Taking the higher percentage, because one would expect

much less recovery at lower temperatures, we estimate that the stored

energy is roughly 0.4 W for the high strain rate and 0.3 W for the lower

strain rate. Therefore, the maximum possible calculated stored energy

(from Eq. 3) is less than the probable experimental value of deformation

work. Thus, either significant work is expended in creation of other

defects, such as dislocation loops, 7 vacancies, or interstitials, or other

forms of the long-range stress fields must be considered. Also, since the

total stored energy is relatively small compared with the work of deformation, considerable heat must be removed from the specimen during a tensile

test

14
H • 1 1: Conduction Through a Solid

The conductance of heat along the specimen, assuming a uniform temperature- gradient across the specimen, is represented as:

Qs - 2KAAT/AX, (5)

with K, the thermal conductivity; A, the cross-sectional area; AT, the temperature change over Ax, the distance of conduction. A factor of two is

included since heat is conducted to both ends of the specimen. At the

strain rate 2.2 x 10' 3 s' 1

, the bulk temperature is assumed to be 4 K, discounting localized thermal spikes that do not affect the ultimate tensile

strength. Thus, Qs for the lower strain rate is zero. At the strain rate

of 44 x 10' 3 s' 1

, the bulk temperature at fracture is assumed to be 150 K,

from the internal temperature measurements of Ogata and Ishikawa. 2 The

temperature gradient is estimated to occur from the fracture surface to the

specimen grips. During strain-rate change experiments, 6 small temperature

increases of less than 1 K were occasionally detected by thermocouples

attached to the grip. Considering that the maximum internal temperature,

measured at 150 K, is associated with specimen necking, Ax is approximately

2.5 x 10' 2 m and the specimen cross-sectional area is 0.32 x 10' m

2

. For

a temperature gradient of 150 K, we assume an average thermal conductivity

for alloy 316LN of 5.5 W-m' 1 -K' 1

. 9 From Eq. 5, one obtains Qq = 2.1 W for

« - 4.4 x 10' 3 s' 1

.

Thus, at the high strain rate (4.4 x 10' 3 s' 1 ), E^ef — Egg — Qs - (5.5 — 0.4 — 2.1) W - 3.0 W, and at the lower strain rate (2.2 x

10' 3 s' 1 ), E^gf — Egg — Qs = (2.9 — 0.3) W = 2.6 W. These excess heat

quantities must be conducted through the specimen surface to liquid He.

Sol id- to-Liquid Heat Conduction

From the above calculation, most of the deformation work is apparently

converted into heat and most of the heat conduction is to the liquid He

rather than through the specimen. Heat conduction to the fluid, Ql, is

Ql ” hK

as at > (6)

where H^ is the heat transfer coefficient at the solid- liquid surface and

As is the surface area. In boiling liquid He at 4 K, H^ represents a bulk

coefficient associated either with steady-state nucleate or film boiling

from the solid surface. The conditions for nucleate or film boiling depend

on the temperature difference and on the heat flux per unit area, Q/As ; these are portrayed in Fig. 15. 10 Thus, when Q/As > 1 x 10 4 W-m" 2 or AT > 10 K, film boiling is present at the specimen surface. The factor

Ql/As for high strain rate is (E<je f — Egg - Qs )/As or 3.0 W/3 x 10' 4 m

2 - 1.0 x 10 W-m' 2

. Since Ql/as is about 1 and AT is approximately 100 K,

Fig. 15 indicates that film boiling is predominant. For the lower strain

rate (E^ef — Egg — Qs )/As =2.6 W/3.0 x 10" 4 m' 2 = 0.87 x 10 4 W-m' 2 and

AT < 1 K; in this regime nucleate boiling is prevalent. Thus, changing

from the strain rate 2.2 x 10" 3 s" 1 to 4.4 x 10' 3 s" 1 results in a change

15
in the heat transfer mechanism to liquid He; nucleate boiling changes to

film boiling. At the higher strain rate, heat transfer cannot keep up with

heat input, and the tensile specimen warms considerably.

Since these calculations depend on the geometry and thermal conductivity of the specimen, they may not be applicable to other specimen shapes

and materials. Alloys with higher thermal conductivities are expected to

remain at 4 K at higher strain rates.

Martensitic Transformation

The divergence of the trends of percent a' martensite as a function of

strain rate (Fig. 14) are, at first, perplexing. The probable explanation

for the increased transformation of alloy 304L and decreased transformation

of alloy 316LN at high strain rates must involve the temperature dependence

of the strain- induced transformation. We have shown earlier5 for highly

metastable austenitic steels, such as 304L, that there is increased straininduced transformation at 76 K and that the transformation is suppressed at

lower temperatures. Temperature measurements of alloy 304L reveal that the

discontinuous yielding events result in instantaneous warm-up to 50 to 100 K,

the temperature range of maximum a’ formation in alloy 304L. In contrast, no suppression of alloy 316LN a' per unit strain is apparent at very low

temperatures. 11 Tobler et al. reported continuously increasing amounts of

a' per unit strain from 175 K to 4 K. Therefore, whereas temperature increases at higher strain rates stimulate a' formation in alloy 304L, temperature increases in alloy 316LN suppress a' formation.

SUMMARY

Tensile tests of alloys 304L, 310, and 316LN were conducted as a function of strain rate in boiling liquid He (4 K) . At strain rates less

than 2.2 x 10' 3 s’ 1

, all tensile properties remain insensitive to strain

rate. At strains rates of 4.4 x 10" 3 s" 1 or larger, ultimate tensile

strengths were less, and stress—strain characteristics were altered.

Measurements indicated specimen warming up to 100 K. From heat balance

calculations, we conclude the following:

1. Relatively little of the work of deformation is retained as stored

energy in the specimen; we estimate about 10%.

2. Heat transfer through the solid specimen is less than that to the

liquid He.

3. Heat transfer to liquid He at strain rates < 2.2 x 10" 3 s' 1 is via

nucleate boiling.

4. Heat transfer to liquid He at strain rates > 4.4 x 10 -3

s' 1 is via

film boiling.

16
Thus, specimen warming during tensile tests in liquid He at high strain

rates (resulting in reduced ultimate strengths) is associated with the

transition from nucleate to film boiling on the specimen surface.

5.

ACKNOWLEDGMENTS

The Office of Fusion Energy (DOE), Victor Der project monitor, sponsored

this research.

The authors thank P. T. Purtscher (NBS ) for the excellent fractographs

and L. A. Delgado (NBS) for assisting in metallography. Lawrence Livermore

National Laboratory provided the 316LN alloy plate.

REFERENCES

1. D. T. Read and R. P. Reed, Adv. Crvo. Eng. —Mater. 26:91—101

( 1980 ).

2 . T. Ogata and K. Ishikawa, Workshop on Standardization of Fracture

Toughness Testing of Low Temperature Structural Materials, National

Research Institute for Metals, Tsukuba Labs., Sengen, Sakura-mura,

Niihari-gun, Ibaraki, Japan, (March 1986).

3 . W. G. Dobson and D. L. Johnson, Adv. Crvo. Eng. —Mater . 30: 185—192

(1984).

4. N. Yamagami, Y. Kohsaka, and C. Ouchi , Report to Joint U.S./

Japan Working Group on Structural Materials at Low Temperatures,

Technical Research Center, Nippon Kokan K.K. , Kawasaki, Japan

(October 1986) .

5. R. P. Reed, "Materials at Low Temperatures," American Society for

Metals, Metals Park, Ohio (1983), pp . 295—341.

6. R. P. Reed and R. P. Walsh, 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.

7. D. Ronnpagel and C. Schwink, Acta Metall. 26:319—331 (1978).

8. T. Ungar, H. Mughrabi, D. Ronnpagel, and M. Wilkins, Acta Metall.

32:333-342 (1984).

9. N. J. Simon and R. P. Reed, in: "Materials for Superconducting

Magnet Systems," National Bureau of Standards, Boulder, Colorado

(June 1982), pp. 316—363.

17
10. E. G. Brentari, P. J. Giarratano, and R. V. Smith, "Boiling Heat

Transfer for Oxygen, Nitrogen, Hydrogen, and Helium," NBS Tech. Note

317, National Bureau of Standards, Boulder, Colorado (Sept. 20,

1965).

11. R. L. Tobler, D. H. Beekman, and R. P. Reed, in: "Materials Studies

for Magnetic Fusion Energy Applications at Low Temperatures—-VI ,

NBSIR 83-1690, National Bureau of Standards, Boulder, Colorado

(1983) pp. 87-112.

Table 1. Material Characterization

Alloy Chemical Composition (wt.%)

Hardness

(RB ) Grain

Size (fim)

304L-A 18.4 Cr, 9.7 Ni, 1.4 Mn,

0.02 C, 0.03 N, 0.02 P,

0.01 S, 0.6 Si, balance Fe

72 70

304L-B same as 304L-A 73 72

310 -24.8 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 F

72 161

316LN-A 18.0 Cr, 10.5 Ni, 1.1 Mn,

0.016 C, 0.167 N, 0.02 P,

0.015 S, 0.5 Si, 2.12 Mo,

0.2 Cu, 0.25 Co, balance Fe

77 78

316LN-B 18.0 Cr, 10.5 Ni, 1.1 Mn,

0.016 C, 0.149 N, 0.02 P,

0.015 S, 0.5 Si, 2.12 Mo,

0.2 Cu, 0.25 Co, balance Fe

81 57

18
Fig. 1. Microstructures of alloy 304L-A (a) and 304L-B (b) with similar

grain size and hardness, but slightly varying yield strength.

19
Fig. 2. Microstructure of alloy 316LN-A (a) and 316LN (b) taken

from the same plate, but with varying grain size, hardness,

inclusion density, and yield strength.

20
100 Mm

Fig. 3. Microstructure of alloy 310, partially sensitized

with carbide formation at grain boundaries.

21
6 . 3 mm

(0.25 in)

, 38 mm

(1.5 in)

71 mm

(2.8 in)

Fig. 4. Tensile specimen.

Fig. 5. Dependence of tensile ultimate strength on strain rate at 4 K.

22
1000

« 900 —

2

X*h

o

LUX

H1 C/5 Q

_j

LU>

800

700 —

400

>—

300

316LN-A _Z

316LN-B

^

310

10 -6 io -5 10 -4 10~ 3 STRAIN RATE, S" 1

Fig. 6. Dependence of tensile yield strength on strain rate at 4 K.

Fig. 7. Dependence of reduction of area and tensile elongation

on strain rate at 4 K.

23
Fig. 8. Tensile stress -vs strain curves at various

strain rates for alloy 310 at 4 K.

100

50

Fig. 9. Tensile stress-vs . -strain curves and specimen surface

temperature at various strain rates for alloy 316LN.

24

TEMPERATURE.
STRESS,

MPa

~

STRESS.

MPa

STRAIN (b) STRAIN

(d) STRAIN

Fig. 10. Tensile stress -vs strain curves and specimen surface

temperatures at various strain rates for alloy 304L.

25

TEMPERATURE,

K

TEMPERATURE,

K
(a) 100 Mm (b) 100 Mm

Fig. 11. Fractograph produced by scanning electron

of alloy 304L tensile specimens at strain

4.4 x 10 s'l (a) and 4.4 x 10 -3 .-1 (b)

microsopy

rates of

(SEM)

26
a) 100 ym 0>) 100 ym

Fig. 12. SEM fractograph of alloy 316LN tensile specimens at a strain

rate of 4.4 x 10'^ s"^. Regions of microvoid coalescence (a)

and shear (b) are shown.

27
Fig. 13. SEM fractography_£)f a^.loy 310 sensitized tensile specimens at a strain

rate of 4.4 x 10 s Part (c) illustrates microvoid coalescence on

the grain boundary. Part (d) is polished and etched (oxalic acid) surface.

I 28 1
Strain Rate (s

1 ) Fig. 14 Normalized amount of a' martensite formed in uniform

elongation portion of 304L and 316LN tensile specimens

at 4 K as a function of tensile strain rate.

Fig. 15. Heat flow per surface area as a function of temperature

difference of smooth, solid, metal surface in liquid helium.

29



MODIFICATION OF 316LN-TYPE ALLOYS

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

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

A series of eight austenitic stainless steels was tested

in liquid helium to determine the effect of Mo and Ni variations on the strength and toughness. The Mo content ranged from

0 to 4 wt.%; the Ni content varied from 11 and 14 wt . % . The

microstructure of the alloys depended upon the composition and

annealing temperature. Higher alloy content and lower annealing

temperatures, 1000 to 1050°C, resulted in a nonuniform structure.

The higher temperature, 1150°C, produced a uniform austenitic

structure. The mechanical test results showed that Mo additions

increased the yield strength (to a maximum at 3 wt.%) and decreased the Kj c (J) values, so that there was no improvement in

the strength—toughness relationship. Increasing the Ni content

decreased the strengthening effect of Mo and increased Kp c (J).

This supports earlier work that showed that Ni does improve the

strength—toughness relationship. The strain- induced martensitic

transformation was suppressed by increases in Ni, Mo, or N.

INTRODUCTION

The structures used to support superconducting magnets are fabricated

from austenitic stainless steels. In the first generation of magnets, types

304LN and 316LN were used extensively because they were commercially available and their mechanical properties met the minimum requirements. The

316LN composition is similar to 304LN except that it contains 2 to 3 wt.% Mo

and 10 to 14 wt.% Ni, rather than the 8 to 12 wt.% Ni specified for 304LN.

Future designs for superconducting magnets will require structural alloys

with improved mechanical properties, particularly the yield strength in

relation to the fracture toughness.

The purpose of this work is to determine whether modification of the

existing 316LN alloy composition would significantly improve the yield

strength and fracture toughness at 4 K. Past experience with 304LN and

316LN 1 has shown that adding Mo and increasing Ni improves the yield

strength at 4 K with little change in fracture toughness (thus, a better

31
strength—toughness combination). In this study, we evaluated the effect of

Mo and Ni additions on the strength and fracture toughness as a function of

annealing treatments.

MATERIALS

Eight 30-kg laboratory heats were vacuum- induction melted and cast

under an Ar atmosphere into round ingots, 127 mm in diameter and 203 mm in

length. The compositions are listed in Table I. Sulfur was removed from

the heats by adding 50 g of calcium silicon to the melt before pouring.

The ingots were forged at 1200°C to billets 76 mm by 76 mm in cross

section. Sections of the billets were hot-rolled to a final thickness

of 25 mm. The hot-rolled sections were annealed in an Ar atmosphere at

temperatures between 1000 and 1150°C followed by quenching in an agitated

water tank. The time and temperature at which each specimen was treated

is shown in Table II along with comments about the resulting microstructure.

Figure 1 shows a representative microstructure of those specimens that con- tained an unknown second phase in the austenitic matrix. The unknown phase

was nonmagnetic and typically had a lower Mo content than the matrix. Fig- ure 2 shows an example of the microsegregation found in several specimens.

The second phase was dissolved, but the etch revealed a "ghost" structure

that is an indication of nonuniform composition.

PROCEDURES

The 4-K mechanical properties were determined according to the proposed

standard procedures set forth by the United States -Japan Cooperative Program

Committee on Test Methodology. These new test standards are based on the

existing ASTM E 8 and E 813 standards. 2 ' 3

The mechanical properties are a function of grain size, N content, and

the test variables, Mo and Ni concentration. Therefore, it is useful to

adjust the mechanical properties to reduce the influence of variables other

than those that we wish to analyze. The yield strengths measured in this

study were adjusted to a nominal N content of 0.20 wt . % and grain size of

50 /xm according to equation 1

:

4

cjy = constant + 2778[N] + 811d~ (1)

where Oy is the tensile yield strength in MPa, [N] is in wt.%, and d is the

austenite grain size in /zm. The constant was calculated for each Oy measured, and then the nominal [N] and d were added to the constant to determine the adjusted Oy.

The fracture toughness is a strong function of the tensile yield

strength. To evaluate the effect of the composition variables on both the

yield strength and fracture toughness, the quality index (QI) , the product

of the yield strength and fracture toughness, can be calculated. The QI

32
gives a direct means of relating metallurgical parameters to the strength—

:uu;;hni'ss relationship. 5 Our previous study showed that increasing the Ni

content in a series of 304LN-type alloys did not affect the yield strength

significantly, but increased the fracture toughness (the QI increased by

approximately 20 GPa-MPa/m); increasing the N content increased the yield

strength and decreased the fracture toughness (no clear effect on the QI)

An adjusted QI (for constant Ni content) was calculated according to the

(.rend previously established, so that any effect of Mo on the strengthtoughness relationship would be shown.

RESULTS AND DISCUSSION

The measured mechanical properties are summarized in Table III. Fracture toughness as a function of yield strength is plotted in figure 3. The

reference trend line refers to a series of annealed Fe-18Cr-10Ni steels with

various C and N contents to increase the strengths. 6 Most data fall within

the range of values reported previously for the reference trend.

From figure 3, we see that the temperature for the annealing treatment

affected the s trength—toughness relationship. Higher annealing temperatures

raised the values (for alloys 0/11, 4/12, and 4/14) into the range previously measured for the reference trend. Metallography of the annealed

specimens revealed the presence of an unknown second phase (see figure 1) in

the microstructure of most of those specimens with a low strength—toughness

relationship. Annealing at higher temperatures dissolved the second phase

and resulted in better properties. For alloys with no apparent segregation

problems, higher annealing temperatures resulted in larger austenite grain

sizes, which produced lower yield strengths and correspondingly higher

fracture toughness values.

To evaluate the effect of Mo on the yield strength, only specimens

in which the second phase was dissolved were considered (elongation > 30%

and reduction in area > 45%) . The adjusted yield strength plotted as a function of Mo content are shown in figure 4. Between 0 and 3 wt . % Mo, the

strength increases, but for the 4 wt.% Mo alloys, the strength levels off.

Two lines are drawn on the figure to reflect a possible influence of Ni on

the strengthening effect of Mo.

To determine whether Ni influences the Mo strengthening, we plotted the

adjusted yield strength as a function of Ni content in figure 5. At a con- stant Ni content, 3 wt.% Mo apparently produces the highest strength, the

same conclusion drawn from figure 4. Also, at the three different Mo

contents, increasing the Ni content decreased the strength.

Next, we considered the effect of Mo on the strength—toughness relationship. The adjusted QI values for specimens with a uniform microstructure are plotted as a function of Mo in figure 6. The adjusted QI does not

increase, as it does for increasing Ni content, but decreases slightly as Mo

is added. Apparently the higher Ni content in 316LN results in properties

superior to those of 304LN.

33
There is a considerable range of QI values for the alloys with about

2 wt.% Mo (316LN) . An earlier report 4 summarizes the properties of 316LN

alloys from different sources. If the QI is calculated and adjusted to

12 wt.% Ni, there is an even larger range of values - -approximately 170 to

270 on the scale of figure 6. Clearly, the QI for 316LN alloys is not a constant, but varies considerably depending upon the heat treatment and

inclusion content of the particular alloy.

The austenitic matrix partially transforms to the a' martensite phase

in these alloys during plastic deformation. The extent of the transformation depends upon the alloy content, the amount of plastic deformation, and

the temperature. The average martensite content after fracture of selected

tensile specimens is divided by the percent elongation for that specimen and

plotted in figure 7 as a function of the Mp, the strain- induced martensite

start temperature, as calculated from equation 2: 7

Md = 686 - 6 [Cr ] - 25 [Ni] - 16[Mn] + 21 [Si] (2) - 222 [C+N] - 11 [Mo]

,

where Mp is in degrees K and the alloy contents are in wt . % . As predicted

by equation 2, the results show that Mo does reduce the Mq

.

CONCLUSIONS

1. Molybdenum additions increase the yield strength at 4 K, up to a maximum at 3 wt.% Mo. The fracture toughness decreases and the

adjusted QI decreases slightly over the range of 0 to 4 wt.% Mo.

2. Molybdenum additions are less effective in increasing the yield

strength of alloys with high nickel content. Nickel and molybdenum

apparently have interactive effects on the yield strength.

3. Higher annealing or processing temperatures are required to produce

a completely austenitic microstructure as the molybdenum content in

the alloy increases.

4. Molybdenum additions, as well as nickel and nitrogen additions,

suppress the strain- induced martensitic transformation at 4 K.

5. The 316LN composition has a better combination of strength and

toughness than 304LN at 4 K because of its higher nickel content.

REFERENCES

1. D. T. Read, and R. P. Reed, Cryogenics . vol. 21, 1981, p. 415.

2. ASTM E 8-85b, in: 1986 Annual Book of ASTM Standards . American Society

for Testing and Materials, Philadelphia, 1986, p. 124.

3. ASTM E 813-81, in: 1986 Annual Book of ASTM Standards . American Society

for Testing and Materials, Philadelphia, 1986, p. 768.

34
4. R. P. Reed, P. T. Purtscher, N. J. Simon, and R. L. Tobler, in:

M itiTi.ils Studies for Magnetic Fusion Energy Applications at Low

Temperature—IX . R. P. Reed, ed. , National Bureau of Standards, Boulder,

Colorado, 1986, p. 15.

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

Cryogenic Engineering , vol. 30, R. P. Reed and A. F. Clark, eds., Plenum

Press, New York, 1986, p. 43.

6. R. L. Tobler, D. T. Reed, and R. P. Reed, in: Fracture Mechanics.

Thirteenth Conference . ASTM STP 743, R. Roberts, ed. , American Society

for Testing and Materials, Philadelphia, 1981, p. 350.

7. I. Williams, R. G. Williams, and R. C. Capellaro, in: Proceedings of

Sixth X nternational Cryogenic Engineering Conference . IPC Science and

Technology Press, Guildford, Surrey, England, 1976, p. 337.

Table I. Chemical Composition (wt.%)

Alloyt Cr Ni Mo Mn N % a'/% elongation Mg(K)

0/11 18.0 10.8 0 1.0 0.13 2.05 261

2/11 18.6

' 10.9 2.2 1.0 0.13 1.88 237

2/12 18 12 2 1.0 0.18 0.75 200

3/12 18 12 3 1.0 0.19 0.75 187

4/12 18.2 12 4 1.0 0.19 0.45 176

2/14 18.5 14 2 1.1 0.17 0.33 152

3/14 18.2 14 3 1.0 0.18 0.12 139

4/14 18 14 4 1.0 0.19 0.08 126

* Md = 686 - 6 [Cr ] - 25 [Ni ] - 16[Mn] - ll[Mo] - 222 [N]

Alloy's designation refers to the Mo content/Ni content

35
Table II. Summary of Heat Treatment

Spec .# Temperature (°C) Time (min) Grain size (/zm) Comments

0/11-1 1000 15 23 - ~ - 0/11-2 1050 60 75 —

0/11-3 1050 1440 80 —

0/11-4 1150 1440 161 —

2/11-1 1000 6 7 —

2/11-2 1000 15 11 —

2/11-3 1000 60 20 —

2/12-1 1000 30 34

2/12-2 1000 30 28 —

2/12-4 1150 120 100 —

3/12-1 1000 30 34

3/12-2 1000 30 35 —

3/12-4 1150 120 70 —

4/12-1 1000 30 25 2nd phase present

4/12-2 1000 30 31 2nd phase present

4/12-4 1150 120 75 —

2/14-1 1000 30 37 micro segregation

2/14-2 1000 30 36 micro segregation

2/14-3 - 1050 60 47 —

2/14-4 1050 240 57 —

2/14-5 1050 1440 127 —

2/14-6 1150 30 68 —

2/14-7 1150 90 94 —

2/14-8 1150 330 118 —

3/14-1 1000 30 29 micro segregation

3/14-2 1000 30 33 micro segregation

3/14-3 1150 120 100 —

4/14-1 1000 30 31 2nd phase present

4/14-2 1000 30 39 2nd phase present

4/14-3 1050 240 59 microsegregation

4/14-4 1050 1200 93 microsegregation

4/14-5 1050 5760 177 microsegregation

4/14-6 1150 45 78 —

4/14-7 1150 210 118 ... 4/14-8 1150 600 187 —

36
Table III. Summary of Mechanical Properties

wt

% Mo

wt.% Ni

Yield

Strength

(MPa)

Adj us ted

Yield

Strength

(MPa)

Ultimate

Strength

(MPa)

%Elone

%R. A.

Fracture

Toughness

(Mpa/m)

Quality

Index

(GPa- MPa/m)

Adjusted

Quality

Indexf

0/11-1 784 960 1654 42/55 193 151 176

0/11-2 715 924 1620 42/51 202 144 169

0/11-3 677 953 1588 41/50 212 144 169

0/11-4 720 994 1606 40/48 269 194 219

2/11-1 1056 1087 1733 36/49 146 154 177

2/11-2 935 1092 1686 41/54 177 165 188

2/11-3 885 1042 1618 41/58 187 165 188

2/12-1 1049 1078 1598 43/50 163

173 173 2/12-2 1014 1048 1602 44/55 172

2/12-4 945 1035 1527 45/55 — — —

3/12-1 1109 1114 1620 41/50 157

169 3/12-2 1119 1124 1617 40/47 148 173

3/12-4 1096 1142 1562 52/62 —

4/12-1*

4/12-2*

1222

1181

1203

1162

1580

1569

22/23

21/28

98

103

120 118

4/12-4 1067 1116 1549 44/54 — — —

2/14-1*

2/14-2*

1063

1060

1128

1125

1532

1546

37/49

39/49

188

202 207 170

2/14-3 931 1012 1523 43/56 211 196 159

2/14-4 934 1026 1503 42/53 226 211 174

2/14-5 863 990 1477 43/54 241 208 171

2/14-6 894 995 1514 41/58 201 180 143

2/14-7 903 1018 1490 37/47 213 192 155

2/14-8 905 1029 1473 40/52 210 190 153

3/14-1* 1129 1150 1558 32/53 167

195 3/14-2* 1109 1130 1565 34/48 181

3/14-3 994 1084 1478 39/59 — — —

4/14-1* 1164 1161 1548 25/33 98 114 69

4/14-2* 1161 1158 1534 26/35 — — —

4/14-3* 942 979 1534 39/48 149 140 95

4/14-4* 968 1027 1523 37/45 165 160 115

4/14-5* 915 997 1493 32/44 145 133 88

4/14-6 1013 1064 1530 34/52 193 196 151

4/14-7 1007 1075 1508 41/55 170 171 126

4/14-8 981 1065 1502 40/49 195 191 146

Not fully austenitic; nonuniform composition

f See Procedures section of text

37
Figure 1

.

Microstructure of specimen 4/14-1, etched in

potassium hydroxide electrolyte.

38
Figure 2. Microstructure of specimen 4/14-3, etched in

oxalic acid electrolyte.

39
YIELD STRENGTH, MPa

Figure 3. Fracture toughness versus tensile yield strength.

co

MOLYBDENUM CONTENT, wt%

Figure 4. Adjusted yield strength versus Mo content for specimens

with a uniform austenitic structure.

40
ADJUSTED

QUALITY

INDEX,

GPaMPa

VTrT

Figure 5. Adjusted yield strength versus Ni content for specimens

with a uniform austenitic structure.

Figure 6. Adjusted Quality Index, QI , versus Mo content for specimens

with a uniform austenitic structure.

41



MOLYBDENUM EFFECT ON Fe-Cr-Ni-ALLOY ELASTIC CONSTANTS*

H. M. Ledbetter and S. A. Kim

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

By ultrasonic methods, using polycrystalline specimens,

we determined the elastic constants of six face-centered-cubic

Fe—Cr—Ni alloys, nominally Fe—19Cr—12Ni (at. pet.). In these

alloys, Mo content ranged up to 2.4 at. pet. We found that Mo

lowers the Young and shear moduli, and it raises the Poisson

ratio. Contrary to expectation (because it increases volume),

Mo raises the bulk modulus. Qualitatively, our results show

that Ni raises the bulk modulus and Poisson ratio, but Ni lowers

the Young and shear moduli. (Nickel decreases the alloy's atomic

volume.) Our discussion includes existing models based on 3delectron theory.

INTRODUCTION ’

Because 3d- transition-metal alloys provide the basis for commercial

austenitic stainless steels, their physical properties provide much practical interest. Recent attempts to improve the traditional steels through

chemical-composition changes create many new questions.

Compared to alloys of simple (s-p-electron) metals, or even noble

metals, the theory of 3d-electron metals remains unsettled. Thus, for alloy

improvement, theory provides mainly guidelines for interpretation; it fails

to provide a scientific basis for physical-property prediction.

Because they relate to various solid-state phenomena—interatomic

potentials, dislocation mechanisms, empirical models, macroscopic plastic

deformation, elastic properties rank high among the physical properties.

They contain enormous information. Gilman 1 remarked that "the most important mechanical characteristic of a crystal is its elastic modulus." Rather

than Young modulus, a physical chemist might choose the bulk modulus, or reciprocal compressibility. 2 Many physicists, 3 including Einstein, 4 saw the

bulk modulus as the principal elastic constant for considering solid-state

properties. For studying interatomic -bonding changes, the Poisson ratio

provides a valuable viewpoint

5 » 6

*Intended for publication in Journal of Materials Research

43
This study considers alloying effects in face-centered-cubic Fe—Cr—Ni—

base alloys. Specifically, it considers the changes in the elastic con- stants caused by adding Mo, a 4d body-centered-cubic element. Indeed, as described below, Mo represents a b.c.c. paradigm. Because it represents a strong b.c.c. element and because it possesses a large atomic volume, one expects large elastic-constant changes.

By ultrasonic methods (near 10 MHz) , we determined the complete

elastic-constant set. We report results for four: E = Young modulus

,

G = shear modulus, B = bulk modulus, and v = Poisson ratio.

Previously, for Fe—Cr—Ni alloys, we reported similar studies for

alloying with substitutional Mn7 and interstitial C-plus-N. 8 In both

studies, these elements increased atomic volume and lowered all the elastic

stiffnesses: B, E, and G. As described in detail by Gilman, 9 when atomic

volume increases, we expect lower elastic stiffness. All simple interatomic

potentials predict this inverse relationship.

MEASUREMENTS

Materials

We obtained polycrystalline alloys from the research laboratory of a major manufacturer. Table I shows their chemical compositions. The alloys

were prepared in a vacuum induction furnace, cast into round ingots 13.3 cm

in diameter and 21.8 cm in length, cropped and forged at 1100°C to billets

7.6 cm in cross section, rolled at 1100°C, to 2.5-cm plate, annealed for

30 min at 1000 °C, and water quenched.

Method

We determined sound velocities by a method described elsewhere. 10

Briefly, 1.5 -cm cubes were prepared by grinding so that opposite faces were

flat and parallel within 5 /zm. Quartz piezoelectric crystals with fundamental resonances between 4 and 7 MHz were cemented to the specimens with

phenyl salicylate. An x-cut transducer was used for longitudinal waves and

an ac-cut for transverse waves. Ultrasonic pulses 1 to 2 cycles long were launched into the specimen by electrically exciting the transducer. The

pulses propagated through the specimen, reflected from the opposite face,

and propagated back and forth. The pulse echoes were detected by the transducer and displayed on an oscilloscope equipped with a time delay and a microprocessor for time -interval measurements. The sound velocity was computed by

v = 2i/t. (1)

Here, i denotes specimen length, and t denotes the round-trip transit time.

On the oscilloscope, t was the time between adjacent echoes, usually the

first and second echoes, and within these the time between leading cycles.

Elastic -stiffness constants were computed from the general relationship

44
c - pv 2

. (2)

Here, p denotes mass density. The usual quasi- isotropic elastic constants

are related to the longitudinal and transverse sound velocities, Vjj and vt , by

2 longitudinal modulus = C# - pv^g , (3)

2 shear modulus — G - pvt (4)

bulk modulus - B - Cj> — (4/3)G, (5)

Young modulus - E - 3GB/(Cj> — G) , (6)

Poisson ratio - v - (E/2G) — 1 (7) - (1/2) (C^ - 2G)/(C i - G).

RESULTS

Table II shows the numerical results for p, Cj> , G, B, E, and u . Figure 1 shows these results graphically. Figure 2 shows the relative

elastic-constant changes. Table III gives the least- squares results

obtained for a linear fit: y = a + bx. This table also shows b/a, the

percentage change in elastic constant per percentage change in alloying

element. Here, x denotes the Mo atomic fraction. For comparison, Table III

contains b/a results obtained previously for Mn7 and for C-plus-N 8 alloying

into Fe—Cr—Ni alloys.

DISCUSSION

Since Mo is a 4d rather than a 3d element, its atomic volume exceeds

considerably that of Fe, Cr, Ni, and the Fe—Cr—Ni alloy. (In A

3

, the volumes are Fe = 11.78, Cr = 12.00, Ni = 10.94, Fe-Cr—Ni = 11.59, 11 and Mo = 15.58.) That Mo's atomic volume exceeds the alloy's effective atomic volume

by 34 percent deserves attention. Small volume changes often correspond to

large physical-property changes. For example, consider Fe's bulk modulus.

From 0 to 300 K, it decreases 3.0 percent, while volume increases only 0.6

percent. Ignoring the thermal contribution to elastic softening, we obtain

AB/B _0.0295

AV/V “ 0.00594 ( 8 )

Similarly, for Fe, from pressure studies, 12 we find

AB/B _AB = _

AV/V AP (9)

Thus, we expect the bulk modulus to decrease approximately 5 percent per

percent volume increase.

45
Studies by Ledbetter and Austin 11 show that Mo's excess effective

volume in these alloys equals 52 percent, far exceeding the expected 34

percent. Models based on volumes and atomic elastic constants fail to

explain this large increase. An atomic-volume linear-rule-of -mixture model,

which probably represents an upper bound, predicts, of course, 34 percent.

A more realistic model, say that of Gschneidner and Vineyard, 13 which con- tains both second-order and third-order elastic constants appropriate to

large local strains, predicts 28 percent, too low by approximately one-half.

Because the bulk modulus relates intimately to the atomic volume, we need to consider why Mo fits so poorly in the host f.c.c. Fe—Cr—Ni crystal

structure. Mo prefers a b.c.c. crystal structure. Among the 4d elements,

Mo shows the highest melting point. Outside the metallic orbital, Pauling 14

assigns Mo a valence of 5.0, exactly the number of bonding electrons for

maximum cohesion. The potency of Mo as a b.c.c. stabilizer emerged in

Pfeil's 15 studies on Zr alloys. Altmann et al. 16 point out that the b.c.c.

stabilizing "should be more marked for Mo and W, where we have a peak of

the d weight." Altmann et al. also point out that, despite favorable size

factors, Mo shows practically no solid solubility in Ag. In their theory

of ferromagnetism and band structure in transition metals, Bader et al. 17

focus on the 3d bonding orbitals. They show that the d

3

, or T X g, orbitals

lie along the four [111] directions in a b.c.c. crystal structure; in an f.c.c. structure, they lie along the six [110] directions. From this

b . c . c f . c . c

.

bonding-orbital symmetry mismatch, we expect strong physicalproperty changes, changes exceeding those expected for crystal- fieldsymmetry changes of s-electron or p-electron orbitals.

Results in Fig. 1 show the remarkable feature that, despite increased

atomic volume, the bulk modulus increases with increasing Mo.

Explanation of this phenomenon lies embedded in quantum-mechanical

d-electron theory, which presents one of solid- state-physics ' s most difficult problems. Despite enormous effort—for example Mott, 18 Pauling, 19

Zener, 20 Bader et al., 17 Goodenough

,

2

1

the problem remains unsolved.

Herring22 gave an elementary critique of the state of d-electrons in transition metals. He concluded that elements in the 3d row should resemble the

free atoms and that most of these metals possess itinerant electrons. This

reference contains Herring's much-cited remark on 3d-electron theories; "It

is like mixing a few liquors in various proportions to get a variety of different cocktails." Mainly from the view of d-band theory, Brooks 23 reviewed

the subject. In his review, Brooks disputes ideas based on crystal-field

theory and separation of d- levels into Eg and T 2 g

subbands. (A recent suc- cessful tight-binding calculation of iron's elastic constants invokes the Eg

orbitals

.

24

)

A theory by Ducastelle 25 leads to simple relationships for the volume

and the bulk modulus. In the theory, the total energy arises from two

sources; d-band energy calculated by a tight-binding approximation and

Born—Mayer repulsive energy. Although this model is too simplified to ex- plain present results, it displays the principal features necessary for an improved calculation. For the Wigner—Seitz radius, rQ , Ducastelle obtained

46
1 p - q

( 10 ) 1 in pC

r ° p - q qAW0

The bandwidth-energy approximation is

ub - U0e-qr -

The repulsive energy is

Ur - Ce~Pr.

(ID

( 12 )

and

A - z ( 10 - z)/20 . (13)

Here, z denotes the average number of electrons per atom in the d-band, and

U - + U r denotes the total energy. For the bulk modulus, Ducastelle

obtained

fa 2 u] r

2 r o

fa 2 u]

B - VQ W

~

9 ~

V=VQ

UrgJ

= pqrgUc/9

r=r. (14)

Thus, Ducastelle 's model predicts that increased atomic volume (VQ =

4n-rg/3) leads to decreased bulk modulus because Uc represents cohesive

energy per unit volume.

Despite the deficiencies (for present purposes) of Ducastelle' s model,

we can use it_to estimate prQ and qrG

. For this estimate, we need one more

relationship among p, q, and physical properties. We choose the bulkmodulus pressure derivative:

Here, v - V/VQ , and we obtain

SB 6 + (p + q)rG

dP 3

( 16 )

Taking Uc = 58.6 GJ/m3 (the value for iron), B = 160.4 GPa, and 3B/5P =

5. 57, 26 we obtain

prQ - 7.37, qrQ = 3.34. (17)

These results compare favorably with Ducastelle 's estimates for a typical

transition metal: prQ = 9 and qrQ = 3. Despite the different interatomic

pair potential, Ducastelle 's model corresponds closely to the results of a Mie—Grtineisen interatomic potential: 27

U = ar_n + 0r_m

, (18)

47
B = mnUc/9V0 , (19)

and

3B/3P = (m + n + 9)/3. (20)

In considering metallic elements, Furth28 concluded that m = 4 and n = 7 for Fe, Co, Ni, and Cu. Furth' s coefficients lead to prQ = 11.58 and

qrG = 2.42.

For Ducastelle's model, we can get

Ur by imposing mechanical equilibrium:

3U/3r = 0.

This leads to

|ub/ur | - q/p,

| Bb/Br | = (2 + pr0 )/(2 + qr0 )

,

the relative contributions of Ub and

( 21 )

( 22 )

(23)

and

|(3B/3P) b/(3B/3P) r | = (10 + 6pr0 + p

2 r2 )/(10 + 6qr0 + q

2 r2

) . (24)

The band- structure term contributes 31 percent to the energy,

64 percent to B, and 73 percent to 3B/3P. Thus, the band- structure energy

contributes more to the higher order elastic constants; contrary to expectation, for this particular interatomic potential, the repulsive energy

contributes less.

Aside from the complications of 3d electrons, another factor arises

in these alloys: magnetization. Even though the alloys show macroscopic

paramagnetism, disordered local atomic moments may affect the physical

properties. The usual view is that magnetization increases volume and

decreases bulk modulus. 29 In extending Ducastelle's model to allow for

electron correlation, Friedel and Sayers 29 concluded that both correlation

and magnetization increase volume and decrease the bulk modulus.

Using a second-order perturbation treatment of electron correlation,

Friedel and Sayers 30 reconsidered this question. They found that the bulk

modulus may increase with magnetization and increasing volume. They expect

such bulk-modulus increases to occur in elements with a high electrondensity-of -states , transition elements to the right: Fe , Cu, and Ni.

Such bulk-modulus increases are consistent with, or predicted by, other

transition-metal electron- theory models. 31-33

Finally, we note that Fig. 1 shows also the effects of alloying Ni into

Fe—Cr—Ni. Nickel decreases effective atomic volume, raises the bulk modulus

and Poisson ratio, and lowers the Young and shear moduli. Roughly, the Ni

effect on the bulk modulus, (AB/B)/(AV/V) , equals —1.62. Because we failed

to study systematically the Ni effects, we avoid comparisons with models,

48
but we remark that Ni parallels Mo in changing B, E, G, and u . These con- clusions apply only to the studied alloys, which contain 12 to 14 percent

N i In another study, where the alloys contained as low as 5 percent Ni, we observed other effects, sometimes opposite.

CONCLUSIONS

From this study, there emerge three conclusions:

1. Added to f.c.c. Fe-Cr-Ni alloys, Mo increases the bulk modulus, B,

and the Poisson ratio, u . It decreases the Young and shear moduli, E and G.

Because Mo increases the effective atomic volume, V, one expects a decreased

bulk modulus.

2. Most existing 3d-electron- theory models predict that B decreases

when V increases; but several models allow, or predict, the reverse effect,

which we observe when we alloy Mo into Fe—Cr—Ni.

3. Nickel alloyed into f.c.c. Fe—Cr—Ni parallels Mo in changing the

elastic constants: B and v increase, E and G decrease. Unlike Mo, Ni

decreases atomic volume, but the decrease is small.

ACKNOWLEDGMENT

Partial support for this study came from the DoE Office of Fusion

Energy. M. W. Austin assisted with measurements.

REFERENCES

1. J. J. Gilman, Austr. J . Phys . 13, 327 (1960).

2. E. A. Moelwyn—Hughes , "Physical Chemistry" (Pergamon, Oxford, 1961),

p. 210.

3. F. Seitz, "The Modern Theory of Solids" (McGraw-Hill, New York, 1940),

p. 373.

4. A. Einstein, Ann. Phvs . 34, 170 (1910—11).

5. W. Koster, Metall. Rev . 6, 1 (1961).

6. H. M. Ledbetter, S. Chevacharoenkul , and R. F. Davis, J. Appl, Phvs . 60, 1614 (1986)

7. H. M. Ledbetter, J. Mater. Sci . 20, 2923 (1985).

8. H. M. Ledbetter and M. W. Austin, Mater. Sci. Eng . 70, 143 (1985).

9. J. J. Gilman, "Micromechanics of Flow in Solids" (McGraw-Hill,

New York, 1969), p. 29.

10. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl, Phvs . 51,

305 (1980)

11. H. M. Ledbetter and M. W. Austin, unpublished research, National Bureau

of Standards, Boulder, Colorado (1986).

12. M. W. Guinan and D. N. Beshers, J. Phvs. Chem. Solids 29, 541 (1968).

13. K. A. Gschneidner and G. H. Vineyard, J. Appl, Phvs . 33, 3444 (1962).

49
14. L. Pauling, "The Nature of the Chemical Bond" (Cornell U.P., Ithaca,

1969), p. 403.

15. P. C. L. Pfeil , AERE Harwell Report M/TN11 (1952).

16. S. L. Altmann, C. A. Coulson, and W. Hume—Rothery , Proc . Roy. Soc.

Lond . A240, 145 (1957) . 17. F. Bader, K. Ganzhorn, and U. Dehlinger, Z . Phys . 137, 190 (1954).

18. N. F. Mott, Adv. Phys . 13, 325 (1964).

19. L. Pauling, Phvs . Rev . 54, 899 (1938).

20. C. Zener, Phvs. Rev . 81, 440 (1951); 82, 403 (1951); 83., 299 (1951);

85, 324 (1952).

21. J. B. Goodenough, Phvs . Rev . 120, 67 (1960).

22. C. Herring, J. Appl. Phvs . 31, 3S (1960).

23. H. Brooks, in "Electronic Structure and Alloy Chemistry of the Transition Elements" (Interscience, New York, 1963), p. 3.

24. J. Hasegawa, M. W. Finnis, and D. G. Pettifor, J. Phys. F: Met. Phys . 15, 19 (1985).

25. F. Ducastelle, J. Physique 31, 1055 (1970).

26. D. Gerlich and S. Hart, J. Appl. Phvs . 55, 880 (1984).

27. E. Gruneisen, in "Handbuch der Physik," Volume 10 (Springer, Berlin,

1926), p. 1.

28. R. Furth, Proc. Roy. Soc . A183, 87 (1944).

29. J. Friedel and C. M. Sayers, J. Phvs. Lett . 38, L263 (1977).

30. J. Friedel and C. M. Sayers, J. Phvs. Lett . 39, L59 (1977).

31. M. Shimizu, J. Phvs. Soc. Japan 44, 792 (1978).

32. G. Hausch, J. Phvs. F: Met. Phys . 7, L127 (1977).

33. J. Kollar and G. Solt, in "Phonons" (Flammarion Sciences, Paris, 1971),

p. 331.

TABLE I. Chemical compositions (mass pet.).

Alloy Mo Cr Ni Mn C N S Si P

1 2.05 18.5 13.8 1.05 0.021 0.180 0.004 0.30 0.002

2 2.06 18.0 12.0 1.00 0.024 0.188 0.004 0.30 0.003

3 2.98 18.2 14.0 1.02 0.022 0.188 0.005 0.30 0.002

4 3.00 18.0 11.8 1.01 0.023 0.200 0.005 0.29 0.002

5 4.02 17.9 14.2 1.00 0.022 0.196 0.004 0.30 0.002

6 4.08 18.2 12.1 1.03 0.020 0.204 0.005 0.28 0.002

50
TABLE II. Mass density and elastic constants.

Alloy Density

(g/cm3

) Ci

(GPa)

G

(GPa)

B

(GPa)

E

(GPa)

V

1 7.972 262.6 77.28 159.6 199.6 0.292

2 7.961 262.9 77.52 159.6 200.2 0.291

3 7.989 262.8 76.72 160.5 198.5 0.294

4 7.980 263.0 77.18 160.1 199.5 0.292

5 8.004 263.1 76.27 161.4 197.7 0.296

6 7.995 263.0 76.44 161.0 198.0 0.295

TABLE III. Elastic -constant

relationship, where x denotes

Except v , units are GPa.

compositional variation based on y = a + bx

atomic fraction. Based on lower Ni alloys

a b b/a

Molybdenum

G 78.7 -93 1 h-

OO B 158.1 -125 0.79

E 202.5 -185 -0.91

i/ 0.287 0.349 1.22

Manganese

a

G 78.3 -21 -0.27

B 158.2 -72 -0.45

E 201.7 -60 -0.30

V 0.288 - 0.035 1o rH CM

Carbon -plus -Nitrogen

b

G 78.4 -59 1o

B 159.6 -147 -0.93

E 202.0 -158 0.78

V 0.288 - 0.034 1O

t-H CM

a Ref. 7.

b Ref. 8.

51
263.3

CO 0.O'w* 262.9

cT

262.5

201

CO a.o 199

HI 197

162

CO £Lo 161

m

159

78

CO 0.wO 77

o

76

1—CO co E

z £

8.010

7.985

UJ tj) a

7.960

o 0.296

h*<

OC 0.294

Z

o

CO

CO 0.292

o

Cl 0.290

1.2 1.6 2.0 2.4

MOLYBDENUM (at. %)

Fig. 1. Compositional variation of p = mass density, C = longitudinal

modulus , G = shear modulus , B = bulk modulus , E = Young modulus

,

and v = Poisson ratio. Solid lines represent higher Ni content;

dashed lines, lower Ni content.

52
Fig. 2. Relative compositional variation of various elastic constants.

The results represent the lower Ni alloys shown in Fig. 1.

53



MOLYBDENUM EFFECT ON VOLUME IN Fe-Cr-Ni ALLOYS*

H. M. Ledbetter and M. W. Austin

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards, Boulder, Colorado

By x-ray diffraction on powder specimens, we determined the

unit-cell size for six face-centered-cubic Fe—Cr—Ni alloys, nominally Fe-19Cr-12Ni (at.%). In these alloys, the Mo content ranged

up to 2.4 at.%. We found that Mo increases volume: 0.45% per at.%

Usual models based on atomic volumes and elastic compressibilities

fail to explain the large volume increase. We ascribe the discrepancy to changes in interatomic bonding, which we describe in

terms of 3d-electron models.

INTRODUCTION

Among a solid's most important physical properties, atomic volume

ranks foremost. Volume relates directly and inextricably to fundamental

interatomic fdrces. Usual models of cohesion deal principally with three

physical quantities: cohesive energy, bulk modulus, and atomic volume.

Consider, for example, a Mie—Gruneisen interatomic potential:

U(r) - Ar

' n - Br' m . (1)

Here, r denotes interatomic spacing; A and n are coefficients of the

repulsive energy; B and m are coefficients of the attractive energy.

In this model, one obtains for the bulk modulus (reciprocal

compressibility) [1]

B0 = V0

(d 2 U/3V2

) = mnU0/9V0 . (2)

V=V

o Here U denotes internal energy; V, volume; UQ , cohesive energy; and VQ , equilibrium atomic volume. (Do not confuse B and B0 .

More than fifty years ago, the strong effect of atomic size on

solid-state alloy properties, especially solid solubility, was established by Hume—Rothery et al . [2,3]. Even earlier, Wasastjerna [4] and

*Intended for publication in Journal of Materials Science .

55
Goldschmidt [5] realized the importance -of establishing elements' atomic

radii. Pauling [6] found the concept of atomic radius essential for

understanding molecular and crystal structures, which play vital roles in

determining properties. Pauling [7] argued that a resonating-covalentbond concept combined with atomic radii gives qualitative explanations of

many properties: magnetic, heat capacity, hardness, compressibility, and

thermal expansion. Waser and Pauling [8] related compressibilities and

force constants to interatomic distance; they ascribed behavioral deviations

to a change in bond type (electronic structure). Friedel [9] explained the

nonlinear volume dependence of alloys as arising from the constituents' com- pressibility differences. Rudman [10] reviewed the atomic volumes of the

metallic elements. He reported a range of atomic volumes from 8 . 1 A

3 for

Be to 116.0 A

3 for Cs. The 3d transition metals from Cr through Ni show a narrow range between 10.9 and 12.2 A

3

.

The atomic-volume and elastic-constant interrelationship is well established. On this, the Waser and Pauling study was mentioned above. At least

since Fuchs [11], we have known that the bulk modulus that arises from the

free-electron gas varies as r' 5

, where VQ = (4/3)7rr^. Gilman [12] showed

that B varies as r' n

, where n equals 4 for alkali metals, 4 for covalent

tetrahedrally bonded crystals, 8 for fee carbides, 4 for fluorite- type

crystals, 4 for ionic crystals, and 3 for solid rare gases. No simple B—

r

relationship arises for the transition metals. The explanation of why B decreases with increasing r clearly lies at some level more fundamental than

the interatomic potential, which imposes no constraint on [

d

2 U(V) / (SV 2

) ] v=Vo

and VQ

. If the interatomic-potential minimum shifts to another equilibrium

volume, why should its curvature change?

In discussing cohesive forces in metals and alloys, Mott [13] cited

experimental and theoretical reasons why the energy of an atom depends on

the volume available to it rather than on the co-ordination number or its

surroundings. Mott emphasized that atomic volume is more important than

interatomic distance, which depends on co-ordination number.

The present study considers the volume change caused by alloying Mo,

a body-centered-cubic element, into a face-centered-cubic Fe—Cr—Ni alloy.

Fe , Cr, and Ni are 3d transition elements. Mo is a 4d transition element

with much larger atomic volume: V(Mo)/V(Fe) = 15.50 A

3 /11.78 A3 = 1.32.

MEASUREMENTS

Materials

We obtained alloys from the research laboratory of a major manufacturer

of Mo. Table I shows their chemical compositions. Alloys were prepared in

a vacuum induction furnace, cast into round ingots 13.3 cm in diameter and

21.8 cm long, cropped and forged at 1100°C to billets 7.6 cm in cross section, rolled at 1100°C to 2.5-cm plate, annealed for 30 min at 1000°C, and

water quenched. To prepare materials for x-ray diffraction, we filed the

plates, sieved the powders to 0.053 to 0.150 mm, and annealed the powders

in evacuated quartz ampules at 1050°C for 35 min.

56
We prepared specimens by pressing the powders into viscous silicone

grease in a flat sample holder. Specimens were placed in a commercial,

horizontally polarized, Bragg—Brentano x-ray diffractometer. A 0.3-mm

(0.18-degree) receiving slit at the goniometer radius of 22.5 cm was stepscanned in 0.02 -degree 29 increments at 60 s per step. The radiation

consisted of Cu K/3 (1,3) , wavelength 1.39225 A, excited at 44 kV and 40 mA.

Using a Ge solid-state detector, we measured each specimen twice. During

measurements, temperature varied between 307 and 310 K.

We determined peak positions by fitting to a Lorentzian:

I - I max 1 +

Vax)

2W2

-2

Here, I max denotes maximum intensity; 9 , diffraction angle; and W, full peak

width at half -maximum intensity. We obtained nine f.c.c. diffraction lines:

111, 200, 220, 311, 222, 400, 331, 420, 422. To determine the unit-cell

size, we used the 422 diffraction line calibrated against silicon (NBS

Standard Reference Material SRM-740A) . For Cu K/3 radiation, the 422 line

appears at 143 degrees in 29 .

RESULTS

Figure 1 shows the principal results: measured unit-cell size, a, together with a linear- least- squares best-fit line (in angstroms):

a = a x + be = 3.5926 + 0.5385c. (3)

Here denotes the zero—Mo intercept; b, the slope da/dc; and c, the atomic

fraction of Mo. For purposes below, we shall rewrite Eq . 3:

a — a x Aa b

= = — = 0.1499

C3 ^ C cl ^

3. ^

(4)

And, using simple models, we shall try to estimate this quantity. For cellsize measurements, we estimate the systematic uncertainty as 0.0006 A and

the random uncertainty as 0.0001 A. For the slope da/dc, we estimate the

uncertainty as 0.05 A. From this result, it follows that the volume change

equals

V - V x cVj = 0.4504 (5)
Thus, volume increases approximately one-half percent per atomic percent Mo.

If volume followed a linear rule-of -mixture , this result implies that Mo's

effective atomic volume exceeds the average Fe—Cr—Ni—alloy atomic volume by

45%.

In Table I, one sees two sets of alloys differing in Ni content by

approximately 2%. The measurement results in Fig. 1 are different for lowNi and high-Ni alloys. Compared with Mo, the Ni effect on volume is negligible, and we ignore it. Thus, the line in Fig. 1 represents the average

Ni content: 13.0 wt . %

.

DISCUSSION

First, we shall consider whether existing simple models explain the

volume change shown in Fig. 1. We consider six models, use them to calculate (d — d 1 )/cd 1 , and show the results in Table II. By d, we mean the

nearest-neighbor spacing. Subscript 1 denotes the Mo-free Fe—Cr—Ni alloy,

the reference state. Other measures of interatomic spacing include the

unit-cell size, a, where

d = ( 72 / 2 ) a, (6)

and the effective atomic volume,

V = (47r/3)r 3 = a

3 /4. (7)

Among these, we find the relationships

Ad Aa Ar

Some of the models calculate AV/V rather than Ad/d. To compare these, we used the relationship

r > 2 AV d

3 - d? d — d x + 3 d — d x +

d — d : V, df - 3 di

fc J di

Vegard's model

In 1921, Vegard [14] proposed that the unit-cell size follows a linear

rule-of -mixture

:

d = (1 - c)d 1 + cd2 . (10)

Here, subscript 1 denotes the solvent, subscript 2 denotes the solute, and c denotes the atomic fraction. Thus,

58
Ad* d 2 d x

( 11 ) cd x d!

Volume model

A simple alternative to Eq . (10) substitutes atomic volume for interatomic distance:

V - (1 - c)V x + cV 2

. (12)

Zen [15] adopted this model. And, in discussing cohesion in metals and

alloys, Mott [13] argued that atomic volume rather than interatomic distance

tends to remain constant. Thus,

Ad 1 AV V 2 - V x cdj 3 cVj 3cV !

‘ v '

Sphere- in-hole model

The above two models focus on atomic size and neglect atomic com- pressibility, or the elastic-stiffness constants. An isotropic material

possesses two independent elastic constants, often chosen as bulk modulus

(reciprocal compressibility), B, and shear modulus, G. Related to these

is the Poisson ratio

1 3B ~ 2

2 3B + G

(14)

The sphere - in-hole model originated with Bitter [16] and received much

elaboration by Eshelby [17]. For the key equations, we take those given

by Teodosiu [ 18 ]

AV

1 + ^G^B,)

1 + (4G 1 /3B 2 ) (15)

Here, AV' represents the solvent-atom—solute-atom volume difference:

AV' = (4tt/ 3) (r| - rf) . (16)

4.4 Friedel model

Based on a modified Eshelby model, Friedel [9] gives the basic

relationship

d — d x d2 — d x x + 1 cdj d x + x + B 1 /B 2

* (17)

If v denotes the Poisson ratio defined in Eq. (14), then

x = (1 + v)/[2(l - 2u )

]

. (18)

59
4.5 Zener's model

Both the sphere- in-hole model and the Friedel model depend on atomic

volumes and on elastic constants, the second-order elastic constants B

and G. The large strains suggest that third-order elastic constants may

enter the problem. In 1942, Zener [19] developed a model that contains one third-order elastic constant: dG/dP, the shear-modulus pressure derivative.

Zener used an atomic-volume ruj e- of -mixture , Eq . (12), and found two volumechange components

:

AV = AV^ 1 ) + AV^ 2 ) . (19)

The two components are given by

AvC 1 ) = 127rr 2 Ar (1 - is x )/(l + is x ) , and

AV( 2 ) 2Ar f dG

!

r i dP

AV ( 1 ) .

( 20 )

(21)

4.6 Gschneidner—Vineyard model

Another model containing a third-order elastic constant arose in a study by Gschneidner and Vineyard [20]. They found that

AV = c(l — c) 87rr(Ar) 2

. (22)

Here

,

r = (1 — c)r x + cr 2. (23)

Note that this relationship contains a c

3 term. For dilute concentrations,

c « 1, Gschneidner and Vineyard found

d — d x d 2 d x

T—

(

o

V. or Ad

cd x di + 2 lap

' bJ kJ (24)

Table II shows results from the above six models. As input for the

calculations, we used r x = 2.5389 A, r 2 = 2.7252 A, B : = 158 GPa, B 2 =

260 GPa, is x = 0.275, G x = 83.5 GPa, G 2 = 260 GPa, and dG/dP = 1.75. The

second-order elastic constants for the Mo-free state were reported by

Ledbetter and Kim [21] and dG/dP by Gerlich and Hart [22]. We took the

remaining input from standard sources.

The results in Table II, ranging from 0.073 to 0.099, fall far below

the observed value: 0.15. For Mo, if we adopt the Pauling radius for co- ordination number twelve [23], 2.792 A replacing 2.7252 A, the model calculations improve, but they remain well below observation.

The measurement-model discrepancy suggests that alloying Mo into

Fe—Cr—Ni involves electron- structure changes. This observation contains no

60
surprise for several reasons. First, we are placing an element that prefers

the b.c.c. structure into an f.c.c. structure. From a chemical -bond viewpoint, an electron cloud with lobes in {111} directions must adjust to electron clouds with lobes in {110} directions [24]. Second, Mo represents no ordinary b.c.c. metal, but a paradigm. Among the 4d elements, Mo shows the

highest melting point. Outside the metallic orbital, Pauling [23] assigns

Mo a valence of 3.0, exactly the number of bonding electrons for maximum

cohesion. The potency of Mo as a b.c.c. stabilizer emerged in Pfeil's [25]

studies on Zr alloys. Altmann et al. [26] point out that the b.c.c. stabilizing effect "should be more marked for Mo and W, where we have a peak of

the d weight." Altmann et al . also point out that, despite favorable size

factors, Mo shows practically no solid solubility in Ag, and Cr, none in Cu.

Thus, when alloyed into an f.c.c. Fe-Cr-Ni alloy, we expect Mo to cause

strong physical-property changes.

In trying to understand the physical properties of these alloys, we confront one of solid state physics's most difficult problems: quantuminechanical theory of d electrons. Despite enormous effort—for example Mott

[27], Pauling [28], Zener [29], Bader et al. [24], Goodenough [30], Friedel

et al. [31] —the problem remains unsolved. Herring [32] gave an elementary

critique of the state of d electrons in transition metals. He concluded

that elements in the 3d row should resemble the free atoms and that most of

these metals possess itinerant electrons. This reference contains Herring's

much-cited remark on the electronic theories: "It is like mixing a few

liquors in various proportions to get a variety of different cocktails."

Mainly from the viewpoint of d-band theory, Brooks [33] reviewed the subject. In his review, Brooks disputes ideas based on crystal - field theory

and separation of d levels into Eg and T : g

subbands. (However, a recent,

successful, tight-binding calculation of iron's elastic constants invokes

the Eg orbitals [34].)

In general, we expect that whatever increases volume will decrease

cohesion and decrease elastic stiffness. The present alloy system violates

this precept. Ledbetter and Kim [21] showed that alloying Mo into Fe—Cr—Ni

increases the bulk modulus. Thus, to understand these alloys, we must identify an electronic effect that simultaneously increases volume and increases

bulk modulus, the VQ and B 0 in Eq . (2).

A theory by Ducastelle [35] leads to simple relationships for the

volume and the bulk modulus. He assumed that the total energy arises from

two sources: d-band energy calculated by a tight-binding approximation and

Born—Mayer repulsive energy. Although this model is too simplified to ex- plain present results, it displays the principal features necessary for an improved calculation. For the Wigner—Seitz radius, r Q , Ducastelle obtained

ro in-B.SqAW0‘ (25)

The bandwidth approximation is

W = W0

e

_cl r

. (26)

61
The repulsive energy is

Er = Ce'Pr (27)

and

A = z ( 10 - z)/20 . (28)

Here, p and q are constants and z denotes the average number of electrons

per atom in the d-band. For the bulk modulus, Ducastelle obtained

w r

2

l o

'3 2 EC

'

3V 2

«. - 9 v=v0

3r 2

i. j

= pq^oEc/ 9 - (29)

Thus, Ducastelle' s model predicts that increased atomic volume (V0 = 47rr£/3)

leads to decreased bulk modulus because E c represents cohesive energy per

unit volume.

Aside from the complications of 3d electrons, another factor arises in

these alloys: magnetization. Even though the alloys show macroscopic paramagnetism, disordered local atomic moments may affect the physical properties. The usual view is that magnetization increases volume and decreases

bulk modulus [36,37]. However, several recent studies [38—41] show that the

bulk modulus may increase with magnetization and increasing volume. On the

basis of their model, Friedel and Sayers [40] expect such bulk-modulus

increases to occur in elements with a high electron-density-of-states

,

transition elements to the right: Fe , Co, and Ni.

Finally, -we note that the measurements shown in Fig. 1 suggest that Ni

content affects the slope da/dc. Lower Ni content gives a lower slope.

This effect presents little surprise. Nickel, with ten electrons outside

the Ar shell, represents the transition metal that differs most from Mo

(with six outside electrons) . Along with Cu, Ni represents a prototype

f.c.c. metal. Thus, we expect its reluctance to accept Mo as an alloying

element

.

SUMMARY

By x-ray diffraction, we measured the volume change accompanying Mo

alloying into an Fe—Cr—Ni alloy. In volume, Mo atoms are approximately 30%

larger than Fe atoms. We found that Mo increases volume by 45%. Existing

models based on atomic volume and elastic compressibility fail to explain

the large volume increase. We ascribe this large increase to electronic

effects associated with substituting Mo, which prefers a b.c.c. crystal

structure, into an f.c.c. crystal structure. Another study showed that the

alloy's bulk modulus increases with increasing volume. This unexpected

result—increased volume, increased bulk modulus—is consistent with some recent theoretical studies.

62
REFERENCES

1. E. Griineisen, in "Handbuch der Physik, Volume 10" (Springer, Berlin,

1926) p.l.

2. W. Hume-Rothery , G. W. Mabbott, and K. M. Channel-Evans , Philos. Trans . Roy. Soc . A233 (1934) 1.

3. W. Hume-Rothery, R. E. Smallman, and C. W. Haworth, "The Structure of

Metals and Alloys" (Inst. Metals, London, 1969) pp . 71—109.

4. J. A. Wasastjerna, Soc. Sci. Fenn. Comm. Phvs . Math . 38 (1923) 1.

5. V. M. Goldschmidt, Ber. Deut. Chem. Ges . 60 (1927) 1263.

6. L. Pauling, "The Nature of the Chemical Bond" (Cornell U.P., Ithaca,

1960) esp. pp . 393—448.

7. L. Pauling, Proc . Roy. Soc. Lond . A196 (1949) 343.

8. J. Waser and L. Pauling, J. Chem. Phys . 18 (1950) 747.

9. J. Friedel, Philos. Mag . 46 (1955) 514.

10. P. S. Rudman, Trans. Metall. Soc, AIME 233 (1965) 864.

11. K. Fuchs, Proc. Roy. Soc. Lond . A158 (1935) 585.

12. J. J. Gilman, "Micromechanics of Flow in Solids: (McGraw-Hill, New

York, 1969) pp. 29-41.

13. N. F. Mott, Rep. Prog. Phys . 25 (1962) 218.

14. L. Vegard, Z. Phys . 5 (1921) 17.

15. E. Zen, Amer. Mineralogist 41 (1956) 523.

16. F. Bitter, Phvs . Rev . 37 (1931) 1527.

17. J. D. Eshelby, Solid State Phys . 3 (1956) 79.

18. C. Teodosiu, "Elastic Models of Crystal Defects" (Springer, Berlin,

1982) pp. 287-295.

19. C. Zener, Trans. AIME 147 (1942) 361.

20. K. A. Gschneidner and G. H. Vineyard, J. Appl. Phys . 33 (1962) 3444.

21. H. M. Ledbetter and S. A. Kim, unpublished research, National Bureau of

Standards, Boulder, Colorado (1986).

22. D. Gerlich and S. Hart, J. Appl. Phys . 55 (1984) 880.

23 . Ref . 6 , p . 403

24. F. Bader, K. Ganzhorn, and U. Dehlinger, Z . Phys . 137 (1954) 190.

25. P. C. L. Pfeil , AERE Harwell Report M/TN 11 (1952).

26. S. L. Altmann, C. A. Coulson, and W. Hume-Rothery, Proc. Roy. Soc . Lond . A240 (1957) 145.

27. N. F. Mott, Adv. Phvs . 13 (1964) 325.

28. L. Pauling, Phys. Rev . 54 (1938) 899—904.

29. C. Zener, Phys. Rev . 81 (1951) 440; 82 (1951) 403; 83 (1951) 299; 85

(1952) 324.

30. J. B. Goodenough, Phvs. Rev . 120 (1960) 67.

31. J. Friedel, G. Leman, and S. Olszewski, J . App 1 . Phys . 32 (1961) 3255.

32. C. Herring, J. Appl. Phys . 31 (1960) 3S.

33. H. Brooks, in "Electronic Structure and Alloy Chemistry of the Transition Elements" (Interscience, New York, 1963) pp . 3—28.

34. H. Hasegawa, M. W. Finnis, and D. G. Pettifor, J. Phys. F: Met. Phys . 15 (1985) 19.

35. F. Ducastelle, J. Physique 31 (1970) 1055.

36. J. F. Janak and A. R. Williams, Phys. Rev. B 14 (1976) 4199.

37. J. Friedel and C. M. Sayers, J. Phys. Lett . 38 (1977) L263.

38. G. Hausch, J, Phvs. F: Met. Phvs . 7 (1977) L127.

39. M. Shimizu, J. Phys. Soc. Japan 44 (1978) 792.

63
40. J. Friedel and C. M. Sayers, J, Phvs . Lett . 39 (1977) L59.

41. J. Kollar and G. Solt, in "Phonons" (Flanunarion Sciences, Paris, 1971)

pp. 331-335.

TABLE I Chemical compositions (wt. %)

Alloy Mo Cr Ni Mn C N S Si P

1 2.05 18.5 13.8 1.05 0.021 0.180 0.004 0.30 0.002

2 2.06 18.0 12.0 1.00 0.024 0.188 0.004 0.30 0.003

3 2.98 18.2 14.0 1.02 0.022 0.188 0.005 0.30 0.002

4 3.00 18.0 11.8 1.01 0.023 0.200 0.005 0.29 0.002

5 4.02 17.9 14.2 1.00 0.022 0.196 0.004 0.30 0.002

6 4.08 18.2 12.1 1.03 0.020 0.204 0.005 0.28 0.002

TABLE II Model predictions of (1/a) (da/dc)

Model Prediction

sphere- in-hole 0.069

Vegard 0.073

Gschneidner, Vineyard 0.086

Friedel 0.087

Zener 0.099

Volume 0.104

Observed 0.150

64
Fig. 1. Compositional variation of unit-cell dimension, a. Curve through

measurement points represents a linear- least- squares fit to six

values. Open circles represent lower Ni alloys; open squares

represent higher Ni alloys. Lower curves represent six model

predictions: (1) sphere- in-hole ; (2) Vegard; (3) Gschneidner,

Vineyard; (4) Friedel; (5) Zener; (6) volume.

65
-


LOW -TEMPERATURE MANGANESE CONTRIBUTIONS TO THE ELASTIC CONSTANTS

OF FACE -CENTERED -CUBIC Fe-Cr-Ni STAINLESS STEEL*

H. M. Ledbetter and S. A. Kim

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

By ultrasonic methods, we determined the elastic constants

between 295 and 4 K of nominally Fe—18Cr—8Ni alloys containing

up to six percent Mn. We report five elastic constants: C% = longitudinal modulus, B = bulk modulus, E = Young modulus, G = shear modulus, and v = Poisson ratio. At all temperatures, Mn

lowers all these elastic constants. Except for u, larger re- ductions occur at 4 K than at 295 K. At 4 K, the bulk modulus

decreases more than the shear modulus: approximately 0.54 and

0.30 percent per percent Mn, respectively. Manganese raises the

magnetic- transition temperature, which occurs between 40 and 90 K,

by approximately 9 K per percent Mn. A simple model predicts the

volume increase accompanying Mn alloying. But, a simple model

fails to predict the elastic-constant reductions; this suggests

magnetic interatomic interactions.

INTRODUCTION

Previously [1], we reported the ambient- temperature contribution of

Mn to the elastic constants of face-centered-cubic (f.c.c.) Fe-Cr—Ni alloys.

Nominally, these alloys contained 18 percent Cr and 8 percent Ni (mass

percentage). Our results showed that Mn lowers all the elastic-stiffness

constants and the Poisson ratio. The largest effect occurred in the bulk

modulus

AB

cB

-0.47. ( 1 )

* Intended for Journal of Materials Science.

67
Here c denotes fractional Mn concentration. Thus, the bulk modulus decreases approximately 0.5 percent per percent solute. Because Mn's atomic mass (54.94) differs only slightly from the alloy's effective atomic mass

(55.33), differences between atomic and mass percentages are small.

Here, for the same polycrystalline alloys, we report further studies:

elastic -constant measurements between 295 and 4 K. These studies answer two

questions. First, how does the low- temperature magnetic phase transition

affect the elastic constants? Second, does the Mn- concentration dependence

at 4 K resemble that at 295 K? In studying similar Fe—Cr—Ni alloys alloyed

with interstitial carbon and nitrogen, Ledbetter et al. [2] found strongly

different composition effects at 4 and 295 K.

Previously [1], we described Mn's peculiar atomic properties and their

possible effect on alloy physical properties. These peculiarities include a 58-atom unit cell, a bulk modulus 0.35 that of Fe's (despite similar atomic

volumes), antiferromagnetic electronic interactions, and negative lowtemperature thermal expansivity.

We determined the elastic constants by measuring mass density and the

near-10-MHz ultrasonic velocities.

MEASUREMENTS

Most of the measurement details we reported previously [1]. We

achieved low- temperature measurements to 4 K by procedures described by

Naimon et al . -

[ 3 ]

.

RESULTS

For ten alloys at both 295 and 4 K, Table I shows the following

measured and computed quantities: p = mass density, v^> = longitudinal

sound velocity; vt = transverse sound velocity; Cj> = longitudinal modulus

= pvj) ; G = shear modulus = pvt ; B = bulk modulus = Cj> — 4G/3; E = Young

modulus = 9GB/(3B + G) ; and u = Poisson ratio = [E/(2G)] — 1. For 295 and

4 K, Figs. 1 and 2 show graphs of the measured quantities plotted versus

composition.

For one alloy, Fig. 3 shows the temperature variation of the elastic

constants

.

Figure 4 shows the effect of Mn on the magnetic- transition temperature. We estimated that the transition temperature for both vj? and vt cor- responds to the temperature halfway between the maximum and minimum sound

velocities. Also in Fig. 4, we include results from magnetic-susceptibility

measurements [4]

.

68
DISCUSSION

First, we consider volume effects. Uncorrected for Cr and Ni fluctuations, our mass -density measurements give the volume change:

AV

cV

Ap — - 0.092.

c p

( 2 )

Here, c denotes fractional Mn concentration. Thus, one percent of Mn raises

the effective atomic volume approximately 0.1 percent. Many models exist

for predicting the alloying volume change. The simplest is a linear ruleof -mixture

V - (1 - c)V x + cV2 . (3)

Here, subscript 1 denotes the reference material, the Fe—Cr—Ni alloy, and

subscript 2 denotes the alloy element, Mn. From our mass-density measurements and corroborating x-ray-diffraction measurements [5], we find

equals 11.60 A

3

. The volume V2 corresponding to a Mn atom in an f.c.c.

crystal structure remains uncertain. For the a-Mn crystal structure, V2 equals 12.21 A

3

. For f.c.c. Mn, Pauling [6] estimated V2 to be 11.58 A

3

. Studies on Mn—Cu alloys predict V 2 values of 12.59 [7], 12.67 [8], and

13.02 A

3 [9]. We adopt a V 2 value of 12.59 A3

. Then Eq. (3) predicts

AV

~r - 0

085 . (4)

A more realistic model, which uses both second-order and third-order elastic

constants, wa§ given by Gschneidner and Vineyard [10]:

Ad d2 d x dG G Ad

cd di + 2 [dP

' bJ kJ

Here, d denotes interatomic spacing; G, shear modulus; B, bulk modulus; and

P, pressure. From Gerlich and Hart [11], dG/dP = 1.75, and d relates to

volume according to

d - (72/2)a = (72/2) (V/4) 1 /3 . (6)

Here, a denotes cubic

AV d

3 - d3 v7

" —

df

Substituting into Eq.

unit -cell dimension, and

(5) and using Eq. (7), we obtain

AV

cV

0.091.

Thus, a simple model predicts the observed volume change.

(7)

( 8 )

69
Second, we consider the observed elastic -constant reductions caused by

Mn alloying. As described elsewhere |;i 2 ], a model by Eshelby predicts the

elastic -constant changes arising from volume change caused by dilatation

centers

AB dB 7 ~ 1 dV

(9)

B dP 7 V

and

AG dG B 1 dV

( 10 ) G dP G 7 V

'

Here

,

. 3B + 4G l-i/

7 = (ID

3B = 3

l + i/

’

From Gerlich and Hart [11], dB/dP = 5.57 and dG/dP = 1.75. Thus, on the

basis of the observed volume change, these relationships predict AB/cB = —0.20 and AG/cG = —0.21. These predictions fall well below observation:

—0.54 and —0.30 for B and G, respectively. This discrepancy, especially

for B, means that a simple model based on volume change and second-order

and third-order elastic constants fails to explain observation. We conclude

that other interatomic interactions occur: probably local magnetic interactions, which affect B more than G. As described below, two other observations support the hypothesis of local magnetic interactions. First, Mn

raises the magnetic- transition temperature, stabilizing an ordered magnetic

phase. Second, upon cooling, far above the magnetic transition, B softens

before G. This suggests that as interatomic spacing decreases, magnetic

repulsion interactions intensify and volume increases, lowering B.

Third, we compare the 295-K and 4-K composition effects. Figures 1 and 2 show their similarity, the 4-K changes being slightly larger. (The

Poisson ratio provides an exception: at 4 K, it changes less with com- position.) This similarity suggests that the Mn—Fe—Cr—Ni interactions

do not differ much at temperatures above and below the magnetic- transition

temperature

.

Fourth, we consider the cooling curves, which reproduced reversibly

during heating. At the magnetic transition temperature (confirmed separately by magnetic-susceptibility measurements [4]), all the elastic

stiffnesses decrease and the Poisson ratio increases. Although we did

not measure the volume change associated with the magnetic transition, we suspect that this expansion during cooling may explain most of the elastic

stiffness decrease. Above, we mentioned that the bulk modulus begins to

soften at temperatures well above the magnetic- transition temperature.

This softening we ascribe to increased (by lattice contraction) local

magnetic interactions, interactions that soften B but not G. Below the

magnetic- transition temperature, all the elastic stiffnesses resume usual

temperature behavior, except the Poisson ratio. Instead of decreasing as expected, it tends to increase. Again, we ascribe this unexpected tendency

to increase as showing changes in magnetic interactions, this time below

the magnetic- transition temperature.

70
Fifth, we consider how Mn changes tfhe magnetic- transition temperature.

If Mn only increased volume, we would expect the transition temperature to

decrease, because the interacting species are farther apart, but, we observe

an increase: approximately 9 K per atomic percent Mn. Again, we suspect

that increased local magnetic interactions enhance the existing tendency of

the Fe-Cr-Ni alloy to order magnetically. Whether the dominant interaction

is Fe—Mn, or otherwise, remains uncertain.

CONCLUSIONS

From this study of nominally Fe—18Cr—8Ni alloys, there emerge five

conclusions

1. Elastic-constant changes caused by Mn alloying at 4 K strongly

resemble the changes at 295 K. Except for the Poisson ratio, the 4-K re- sults show slightly larger reductions. Thus, cooling through the magnetic

phase transition produces no large changes in the Mn elastic-constant

contributions

2. As expected from previous studies, cooling through the magnetic

phase transition decreases all the elastic stiffnesses and increases the

Poisson ratio. Above the magnetic phase transition, the bulk modulus begins

to soften.

3. Manganese raises the magnetic- transition temperature by a large

amount: 9 K per atomic percent Mn.

4. A simple model based on atomic volumes and second-order and thirdorder elastie'eonstants explains the volume increase.

5. Simple models fall short by approximately one-half in explaining

the elastic-constant decreases. Especially, the bulk modulus decreases more

than expected. We attribute this softening to local magnetic interatomic

interactions

ACKNOWLEDGMENT

This study received support from the U.S. DoE Office of Fusion Energy.

M. W. Austin consulted on the sound-velocity measurements.

REFERENCES

1. H. M. Ledbetter, J. Mater. Sci . 20 (1985) 2923.

2. H. M. Ledbetter, M. W. Austin, and S. A. Kim, Mater . Sci . Eng .

forthcoming

3. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics 14 (1974)

246.

4. E. R. Jones, T. Datta, C. Almasan, D. Edwards, and H. M. Ledbetter,

Mater . Sci . Eng

. forthcoming.

71
5. H. M. Ledbetter and M. W. Austin, unpublished research, National Bureau

of Standards, Boulder, Colorado (1986).

6. L. Pauling, "The Nature of the Chemical Bond" (Cornell U.P., Ithaca,

1960) p. 403.

7. F. T. Worrell, J. AppI. Phvs . 19 (1948) 929.

8. ASTM card 17-910, American Society for Testing and Materials,

Philadelphia

.

9. P. Venkateswararao and D. K. Chatterjee, J. Mater. Sci . 15 (1980) 139.

10. K. A. Gschneidner and G. H. Vineyard, J. AppI. Phvs . 33 (1962) 3444.

11. D. Gerlich and S. Hart, J. AppI. Phvs . 55 (1984) 880.

12. H. M. Ledbetter and M. W. Austin, Mater. Sci. Eng . 70 (1985) 143.

TABLE I Measured and derived quantities

Mn

(mass %)

P (g/cm3

) (cm/7is)

vt (cm//xs)

Ci

(GPa)

G

(GPa)

B

(GPa)

E

(GPa)

V

T = 295 K

1,02 7.904 0.5752 0.3145 261.5 77.18 157.3 201.2 0.287

2.01 7.861 0.5761 0.3145 260.9 77.75 157.2 200.3 0.288

3.85 7.875 0.5738 0.3143 259.3 77.79 155.6 200.0 0.286

5.81 7.862 0.5717 1.3131 257.0 77.08 154.2 198.2 0.286

T = 4 K

1.02 7.975 0.5783 0.3206 266.7 81.94 157.4 209.5 0.278

2.01 7.931 0.5794 0.3204 266.2 81.39 157.7 208.3 0.280

3.85 7.9448 0.5761 0.3193 263.7 80.98 155.7 207.0 0.278

5.81 7.9329 0.5742 0.3185 261.5 80.45 154.3 205.6 0.278

72
Fig. 1. Compositional variation of several measured quantities:

Cj> = longitudinal modulus, B = bulk modulus, G = shear modulus,

E - Young modulus, v = Poisson ratio. Temperature equals 295 K.

Fig. 2. Companion figure to Figure 1, except temperature equals 4 K.

73
Fig. 3. For the 3.85-Mn ailoy, variation with temperature of several

measured quantities: B = bulk modulus, G = shear modulus,

E = Young modulus, v = Poisson ratio. Nonregular behavior

arises from a reversible magnetic phase transition. The dashed

line indicates temperature corresponding to cusp in magnetic

susceptibility versus temperature. The negative slope di//dT

suggests further low- temperature magnetic changes. Softening

of B during cooling suggests magnetic changes premonitory to

paramagnetic—antiferromagnetic (Neel) transition.

74
TRANSITION

TEMPERATURE

(K)

Fig. 4

Mn (wt. %)

Manganese alloying effect on magnetic (Neel) transition

temperature. Chi denotes dc magnetic susceptibility.

75



Journal of Nuclear Materials 141-143 (1986) 44-48

North-Holland, Amsterdam

STRENGTH AND TOUGHNESS OF AISI 304 AND 316 AT 4 K

*

N.J. SIMON and R.P. REED

Fracture and Deformation Division, National Bureau of Standards, 325 Broadway, Boulder, CO 80303, USA

Structural design of superconducting magnets in fusion energy devices requires reliable property data at 4 K. Nitrogenstrengthened AISI 304 and 316 stainless steels are considered to be the best currently available low temperature structural

alloys on the basis of their fabricability and their potential to meet the US fusion research goals of combined 1000 MPa yield

strength and 200 MPa m

l 2 fracture toughness at 4 K. This paper presents provisional equations for the yield strength of 304- and 3 16- type alloys at 4 K as a function of N content and grain size. Provisional equations for the 4 K fracture toughness.

Kw iJ), 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 senes stainless steels to achieve the strengthtoughness goals of the US 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 the vanous 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. 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 (ov ) by ap- proximately 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 200 /im results in a decrease of o

y 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 energy is dissipated by plastic deformation

during the void growth typical of ductile fracture. We

have shown that the elastic stress intensity factor Klc (J)

(obtained from measurement of the elastic/plastic criti- cal /-integral (/lc ) and calculated from Klc (J) =

(Jlc E) l/2

, where E is Young’s modulus) is linearly

inversely proportional to ov at 4 K [4], Toughness also

* Work supported by the US Department of Energy; not

subject to US copyright.

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 inclu- sion 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 ± 100%. 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

high strength 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 these alloys in a low

temperature structural application, additional heat-by- heat qualification tesing is imperative.

2. Procedure

2.1. Data matrices A series of specimens from both laboratory and

commercial heats of AISI 304- and AISI 316-type alloys were tested. The alloys were prepared in the US, Japan,

Europe, and the USSR [3,4,8], Measurements were made

of fracture toughness [A" lc (/)], 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

0022-31 15/86/S03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

77
46 N.J. Simon, R.P. Reed / Strength and toughness of AISI 304 and 316

30 measurements is:

a

y (MPa) = 215 + 3190[N] + 1093 d~ wl . (4)

The standard deviation (SD) of the fit to the data is 84 MPa; the SDs 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:

av (MPa) = 381 + 2776 [N] + 811<T 1/2

. (5)

The SD of the fit to the data is 55 MPa; the SDs of the three coefficients are 34, 193, and 199. Figs. 1 and 2 show the measured a

y versus the <jy predicted by eqs.

(4) and (5) for both alloys. The dependence of o

y on N in eqs. (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

o0 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-tvpe alloys are

frojn laboratory heats (27 of 30). Perhaps these data

reflect better quality control of the producers in the

larger commercial heats.

Fig. 2. Measured tensile yield strength at 4 K of AISI 316-type

alloys versus tensile yield strength calculated from eq. (5). The

scatterband indicates ±2 SDs. Keys to symbols representing

test specimens are given in ref. [8].

3.2. Fracture toughness

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

KIc (/)(MPa m

1 /2

) = 209 - 0.226

o

y + 14.4 ([Ni] + 0.5 [Mn]) + 294>z~ 1/2

, (6)

where the SD of the fit to the data is 31 MPa m

1 2 and

the SDs of the coefficients are 40, 0.026, 2.9 and 316. For the 316 data, the analogous result is:

AT Ic (/)(MPa m

1 /2 ) = 130 - O.338oy + 20.2 [Ni] + 2252n _1/2

. (7)

The SD of the fit to the data is 31 MPa m

12

, and the SDs of the coefficients are 89, 0.59, 5.2 and 310. Figs. 3 and 4 show the measured Klc (J) versus the Klc (J)

predicted from eqs. (6) and (7). Eqs. (6) and (7) imply

that the fracture toughness of the AISI 316 alloys is more sensitive to all major variables than that of the AISI 304 alloys. The explanation of this awaits more

detailed studies of the effects of metallurgical variables on the /-integral fracture toughness at low temperatures. The SDs of the fits of both alloy types eqs. (2) were

considered to be small enough (31 MPa) that the equations can give useful predictions for alloy design. HowFig. 3. Measured fracture toughness at 4 K of AISI 304-type

alloys versus fracture toughness calculated from eq. (6). The

scatterband indicates ±2 SDs. Keys to symbols representing

test specimens are given in ref. [8].

78
N.J Simon, R.P. Reed / Strength and toughness of AISI 304 and 316 45

test results arc large enough to obscure dependence of cryogenic properties upon parameters.

There were 25 rows in the complete matrix for AISI

)04>type alloys and 19 rows in the complete matrix for AISI 316-type alloys. Additional tensile data without

corresponding Klc (J) measurements were used in the regression analyses of the tensile yield strength. Nitro- gen content ranged from 0.03 to 0.29 wt.% for AISI 304- and from 0 03 to 0.17 wt.% for AISI 316-type alloys.

Nickel content varied from 7.2 to 14.9 wt.% for AISI

304-, and from 10.5 to 13.9 wt.% for AISI 316-type

alloys. Manganese content ranged from 1.0 to 10.22 wt.% for the AISI 304- type alloys. Variation of Mn and

Ni content outside the AISI 304 specifications per- mitted 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 varied from 19 to 385 jim in the AISI 304 matrix,

and from 11 to 138 /im in the AISI 316 matrix. Inclu- sion spacing (defined as n~

x where n = number of

inclusions per mm:

) ranged 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 (/? 0 ) was not used in the

analysis, but was found to vary from 72 to 90 for the AISI 304-tvpe specimens and from 79 to 100 for the AISI 316-type specimens. The complete data matrices

will be published elsewhere [3,8]. The Klc (J) values obtained from the relationship Klc (J) = [/,c £/(l -

i/ : )]' 3

, where v is Poisson’s ratio, were adjusted in accord with the formulation Klc (J) = (JXc E) l/2

. 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

[9], the form of the equation fitted to the data was

0y “ 00 + °n[N] + od d~ x/2

, (1)

where [N] is wt.% N, d is the grain size (/im), and o0

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 re- quire 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 Klc (J) upon yield strength at 4

K has been observed previously in our laboratory [10,11],

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: KXc (J) = k0 + k ao

y + k N ([Ni] + 0.5[Mn]) + kj{n ,

( 2 ) where [Ni] represents wt.% Ni, [Mn] wt.% Mn, and k 0

is a constant. Another way of handling the effect of

inclusion spacing is to to normalize Klc (J) as follows: Kic(J )/n

~ 1/2 = k'0 = k'„o

y + k'N ([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 eqs. (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

PREDICTED TENSILE YIELD STRENGTH, oy

, MPa

Fig. 1. Measured tensile yield strength at 4 K of AISI 304-type

alloys versus tensile yield strength calculated from eq. (4). The

scatterband indicates ±2 SDs. Keys to symbols representing

test specimens are given in ref. [8],

79
N.J. Simon, R.P. Reed / Strength and toughness of AISI 304 and 316 47

PREDICTED FRACTURE TOUGHNESS, K, c

(J), MPa-m V2

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

alloys versus fracture toughness calculated from eq. (7). The

scatterband indicates ±2 SDs. Keys to symbols representing

test specimens are given in ref. [8],

ever, there are significant differences in the coefficients

for AISI 304- and 316-type alloys. In contrast, the coefficients resulting from fits of eq. (3) to both alloy

series were remarkably similar. The SDs of these fits to eq. (3) were higher than for the fits to eq. (2), so these equations are not set forth here but are discussed elsewhere [8]. Compared to a fit of Klc (J) on o

y alone, the additions of both the n~

1/2 term and the [Ni] or ([Ni] + 0.5 [Mn]) term were effective in reducing the SD of the fits, except for eq. (6), in which the addition of the n~

1/2 term does not affect the SD. This term is presented for comparison with eq. (7). It is difficult to measure inclu- sion spacing unambiguously [3], but inclusions do have

a pronounced effect upon the 4-K fracture toughness of

the AISI 316-type alloys. Further inclusion studies should be carried out on commercial heats of AISI

304-type alloys.

3.3. Alloy design

Eqs. (4) through (7) can be used to set provisional

design values for yield strength and fracture toughness. For example, to achieve the US research goal of a

y =

1000 MPa [1] for a 304LHN (316LHN) with [N] = 0.20 wt.%, a grain size of less than 55 (162) /im is required.

(The uncertainty in grain size is ± 25 (44) /2 m, 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 /im is difficult for most mills when rolling thick-section (50-120 mm) plate, this implies that AISI 316- type alloys would be a sounder

choice to achieve such high strength.

Assuming further that [Ni] = 12 wt.% and [Mn] = 2 wt.%, achievement of the US research goal of Klc = 200

MPa m

1/2 would require minimum inclusion spacing of 0.10 (0.073) mm for AISI 304LHN (316LHN). 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 [12],

4. Summary

Matrices of NBS measurements of o

y and Klc (J) at 4 K that include numerical data on alloying, refnining,

and processing variables were analyzed using Unear

regression methods. The AISI 304- and 316-type alloys tested were obtained from a variety of sources and

included both commercial and laboratory heats. Tensile yield strength: Provisional equations were

developed for estimating o

y at 4 K from [N] and d~ x/1

. Fracture toughness: Provisional equations for esti- mating K\C (J) at 4 K from a

y

, [Ni], [Mn], and n~ 1/2 were determined. In addition, the predictions of the quantity Klc ( J)/n~ X/

2

from o

y 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 in placing tentative estimates on at- tainable fracture toughness and tensile yield strength at 4 K for austenitic stainless steels.

The research was partially supported by the Office of Fusion Energy, US Department of Energy, V. Der.

Project Monitor. We are indebted to R.P. Walsh, P.T. Purtscher, and L.L. Scull for many of the measurements. Provision of alloys from a number of sources is described and acknowledged elsewhere [8].

References

[1] H.I. McHenry, in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures- VIII, Internal Rept. 85-3025, Ed. R.P. Reed (US National Bureau of Standards, Boulder, Colorado, 1985) p. 21.

[2] R.P. Reed and N.J. Simon, in: Advances in Cryogenics

Engineering - Matenals, Vol. 30, Eds. A.F. Clark and R.P. Reed (Plenum Press, New York, 1984) p. 127.

80
48 N.J Simon, R.P Reed / Strength and toughness of AISI 304 and 316

( 3 1 R'P Reed. PT Purtscher, N.J Simon, and R.L. Tobler, submitted to Proc. 11th Intern. Cryo Engr. Conf. (But- terworth Scientific, London 1986).

[ 4 ] R L Tobler. D T. Read, and R. P Reed, in: ASTM STP

743 (American Society Testing Materials, Philadelphia,

Pennsylvania, 1982) p. 250.

[5] T Sakamoto, Y Nakagama, I. Yamauchi, and T. Zaizen, Advances in Cryogenic Engineering- Materials, Vol. 30, Eds. A F Gark and R P Reed (Plenum Press, New York,

1984) p. 137.

[6] R P Reed and P T. Purtscher, Advances in Cryogenics

Engineering - Materials, Vol. 32, Eds. R.P. Reed and A.F.

Clark (Plenum Press, New York, 1986) p. 43, in press.

[7] N J Simon and R.P Reed, Preliminary Draft, Structural Materials for Superconducting Magnets -III, AISI 304

Stainless Steel (US National Bureau of Standards, Boulder,

Colorado, 1985), unpublished.

[8] N.J. Simon and R.P. Reed, in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures- IX, Internal Rept., Ed. R.P. Reed (US National Bureau of Standards, Boulder, Colorado, 1986), to be

published.

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

[10] R.L. Tobler and R.P. Reed, in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures- III, Internal Rept. 80-1627, Ed. R.P. Reed, (US

National Bureau of Standards, Boulder, Colorado, 1980)

p. 15. [11] R-L. Tobler, T.A. Siewert and H.I. McHenry, Cryogenics

(1986), to be published.

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

81



PROPERTIES OF CDA 104, 155, AND 175 COPPER ALLOYS

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

National Bureau of Standards

Boulder, Colorado

Three copper alloys, designated 104, 155, and 175 by the

Copper Development Association (CDA) , were evaluated for use as the structural conductor in the compact ignition tokamak project.

Mechanical properties and electrical resistivity were measured

for the annealed, half-hard, and hard conditions of alloys 104

and 155 and for the hard condition of alloy 175. Tensile, creep,

fatigue, and combined creep and fatigue were measured at 4, 76,

and 295 K. Extensive metallurgical characterization of grain

structure, grain size, and hardness was obtained. Effects of

specimen orientation relative to the rolling direction and of

load- and displacement- control tensile tests were studied. Data

are presented in both tables and graphs; metallurgical characteristics are shown in micrographs. Cold-rolling increased

yield strengths, but at the expense of electrical conductivity,

ductility, and toughness. On the basis of our tests, alloy 104

in the hard condition had the best combination of yield strength

and electrical conductivity.

INTRODUCTION

The compact ignition tokamak requires a high-conductivity, highstrength conductor to produce pulsed, high magnetic fields at ambient temperatures. The task is to achieve high strength, good conductivity, and

high material consistency by means of thorough, but inexpensive, alloy

design and processing. Copper and copper alloys are candidate materials.

Three copper alloys were evaluated in this study. To assess the influence of cold- rolling , materials were tested in the half-hard and fully hard

83
conditions as well as the annealed condition. At 4, 76, and 295 K, tensile,

electrical resistivity, fatigue, and creep properties were measured; to

simulate pulsed-magnet operating conditions, a combination of fatigue and

creep was also measured. For some alloys, measurements were taken in both

the longitudinal and transverse orientations. The distinction between

load- and displacement-control testing was also determined, because load

control from magnetic forces is considered the most dominant condition in

magnets. Grain size, hardness, and microstructure were characterized for

each alloy and condition.

This report summarizes all data and experimental techniques; the

results are discussed briefly. More extensive analysis and discussion of

these results will be published later.

MATERIALS

Three copper alloys, designated 104, 155, and 17510 by the Copper

Development Association (CDA) , were included in this study. The CDA 104 was received as 32-mm- thick plate in the half-hard condition; the CDA 155, as 39-mm-thick plate in the half-hard condition; and the CDA 17510 (abbreviated CDA 175), as 25-mm-thick plate in the hard condition.

The hard conditions (H) for alloys 104 and 155 were obtained by coldrolling the as-received half -hard (1/2 H) plate. The annealed conditions (A)

for alloys 104 and 155 were obtained by heat treatments at the National

Bureau of Standards. A summary of the alloys, conditions, and compositions

appears in Table I.

Metallurgical specimens for each alloy and condition listed in Table I were cut from the same plates as the tensile specimens. The specimens were

blocks, with their largest dimension in the longitudinal (rolling) direction;

thus the orthogonal directions of the specimen were maintained throughout the

process of specimen preparation.

After the specimens were mechanically polished with a 220-grit abrasive,

hardness was measured on the surface perpendicular to the short transverse

direction. Then three surfaces of the specimens were mechanically polished

with a 0.05-/im alumina abrasive and chemically etched with an aqueous solution of potassium dichromate and sulfuric acid.-*- The grain structures of the

specimens were observed with an optical microscope; photomicrographs of the

three polished surfaces of the specimens were made (Figs. 1—7). In the

annealed condition, the recrystallized grain structure was fairly equiaxed.

In the half-hard condition, some texture was observed, a result of the coldrolling. In the hard condition, the grain structure of each alloy was quite

lamellar

.

84
EXPERIMENTAL PROCEDURES

Mechanical Property Measurements

The tensile properties of three copper alloys were characterized by

(a) standard displacement-control tensile tests, (b) modified load-control

tensile tests, (c) fatigue tests, (d) creep tests, and (e) fatigue-pluscreep tests. The tests were performed at three temperatures: 295 K (ambient), 76 K (liquid nitrogen), and 4 K (liquid helium).

Tensile specimens taken from the plate material were rods, 6.3 mm in

diameter and 38 mm in gage length, with threaded ends. For rolled plate

materials, the specimens were oriented so that the tensile force was applied

transverse to the rolling direction.

For all tests, specimen strain was monitored with three clip-on straingage extensometers , 25 mm in gage length. They were mounted on the specimen

diametrically opposite and equally spaced over a 25-mm gage section of the

38 -mm specimen gage length. They were excited with a 5 V dc power supply;

their outputs were parallel electronically to give an average of the three

strain measurements. From specimen to specimen, the outputs of these extensometers reproduced within 14 /j. m with a sensitivity of 0.03 strain per volt

(corresponding to a readout of ±0.04 /zm) .

a. Displacement-control tensile tests

The tensile testing machine had a cryostat fixture mounted to the

bottom of its screw-driven crosshead (see Fig. 1). The crosshead displacement rate was -0.2 mm/min. An XY chart recorder recorded specimen load and

the strain measured by the extensometers.

Ultimate tensile strength, 0.2% offset yield strength, elongation,

and reduction in area were determined from displacement-control tests.

b. Modified load-control tensile tests

The programmable servohydraulic test machine was equipped with a cryostat test fixture to enable testing in liquid nitrogen and liquid

helium. The machine was programmed for a ramp function in a load-control

mode; that is, the machine uniformly increased the load on the specimen

until failure. In the previous displacement-control tests, the loading rate

in the linear-elastic region of the stress—strain curves was calculated to

be 57 MPa/min. This loading rate was used in the load-control tests. Thus,

the strain rate of the load-control tests was identical to the displacementcontrol tests until yielding occurred. As soon as the stress—strain curve

deviated from linearity, the specimen strain rate (e) began to increase

(from 0.8 x 10

5 s" 1 ) as the machine tried to maintain a constant loading

rate, and it continued to increase until failure. The maximum strain rate

was actually dictated by the maximum actuator velocity of the test machine.

The maximum strain rate measured during discontinuous yielding was 0.8 s' 1

.

85
Ultimate tensile strength, 0.2% offset yield strength, elongation, and

reduction in area were determined from the load- control tensile tests.

c. Fatigue tests

Fatigue tests were performed to measure the resistance of the copper

alloys to high- stress- level fatigue cycles. The servohydraulic test machine

was programmed for a fatigve cycle with a tension—tension triangular wave- form at a frequency of 0.10 Hz. The peak stress levels of the waveform were

0.9 and 0.1 of the yield strength (cry) . Thus, during the fatigue cycle, the

test specimen was loaded to 0.9cry in 5 s and then unloaded to O.lcry in 5 s

;

this stress spectrum is delineated in Fig. 2. These tests were run

in load control to maintain peak stress levels after plastic strain had

occurred in the specimen. All fatigue tests were run for a minimum of 5000

cycles. Specimen stress and strain were recorded on an XY chart recorder at

chosen places in the fatigue cycle. Measurements were recorded more frequently early in the test when there is a large amount of plastic strain per

cycle

.

d. Creep tests

Because copper conductors could be subjected to high stress levels

throughout the life of high-power magnets, creep tests were performed at

295, 76, and 4 K.

In these tests, the specimen stress was ramped up to 0.9ay in 5 s and

held at this level for a minimum of 1000 min. This was done on a servo- hydraulic test machine in the load-control mode. The load—time history is

shown in Fig. 2. The 25-mm gage-length extensometers monitored specimen

strain within~5 /xe as a function of time. During the test, the stress level on the specimen was maintained within 0.1%.

e. Fatigue-plus-creep tests

In these tests, the specimens were subjected to a combination of

fatigue and creep stresses. To accomplish this, a trapezoidal waveform

(Fig. 2) was programmed on the servohydraulic test machine. During one test

cycle, the specimen tensile stress was increased uniformly up to 0.9ay

in

5 s, held for 10 s, and then decreased uniformly to O.lcry in 5 s and held

for 10 s. Since the machine was in the load-control mode, the peak stress

levels were maintained after plastic strain had occurred in the specimen.

During the 5000 -cycle -minimum test, single-cycle stress—strain measurements

were recorded with an XY chart recorder at chosen places in the fatigue

cycle. Measurements were recorded more frequently early in the test when

there is a large amount of plastic strain per cycle.

Electrical Resistivity Measurements

A conventional, four-probe resistivity technique^ with a commercial

nanovoltmeter was used; the distance between voltage taps was 14.0 cm. The

sample current was in the range 1.9 to 2.0 A for all measurements, giving

voltages in the microvolt range. Potential heating effects were evaluated

86
by measuring the resistance as a function of current at room temperature.

The resistance was linear with current to well above the chosen current.

Resistance measurements were made at room temperature and with the

sample immersed in liquid nitrogen (76 K) and in liquid helium (4 K) . At

each temperature, an automated data acquisition system made twelve sets of

measurements. Each set consisted of readings with the current flow in the

sequence forward—reverse—forward. The program monitored drift of the readings and remeasured the resistance if the variation between the two forward

readings exceeded 0.1%. Precision measurement of the sample cross section

enabled calculation of the resistivity to an accuracy better than 1%.

EXPERIMENTAL RESULTS AND DISCUSSION

Tensile Tests

Tensile properties, including yield strength (cry) , ultimate strength

(au ) , elongation (EL), and reduction of area (R.A.) are given in Tables II

through IV for each alloy and each condition. All data are included. The

strain rate for displacement-control tests and the load rate for loadcontrol tests are identified.

The results for each alloy and condition are plotted individually

against temperature in Figs. 10 through 16. In Figs. 17 through 20, all

data for each tensile property of transversely oriented specimens are com- bined and plotted versus temperature. The stress—strain curves of alloys

104, 155, and 175 under various processing and test conditions are shown in

Figs. 21 through 23, 24 through 29, and 30, respectively.

The ultimate strength was strongly dependent on temperature and only

mildly dependent on cold working and alloying. For the annealed alloys,

the ultimate strength at 4 K was nearly twice that at 295 K (see Fig. 18).

Even for the strengthened alloys, the ultimate strength was at least 20%

higher at 4 K than at 295 K.

Elongation for all alloys and conditions increased at lower temperatures (Fig. 19). Compared with values at room temperature, elongation at

4 K was about 50% higher for annealed alloys and about 100% higher for the

cold-rolled alloys. This disparity was due, in part, to the incursion of

localized necking during room- temperature tests at very low strains (discussed below)

Reduction of area remained nearly constant as a function of temperature (Fig. 20), but there were distinctions between specimen necking at

room temperature and at low temperatures, especially in the cold-rolled

alloys. As illustrated in Figs. 22, 23, 25, and 26, the ultimate load was

reached at relatively low strains (usually <0.02) at room temperature in

both displacement- and load-control tests. Subsequently, during load or displacement extension, the specimen began to neck. As depicted by

these engineering stress—strain curves, when the load began to drop in

87
displacement-control tests (and even in load-control tests) , the pace of

the extension could not keep up with specimen straining; probably adiabatic

heating conditions existed in the necked region. At low temperatures,

the specimens strained more uniformly; consequently, more elongation was achieved, and the reduction of area reflected both uniform elongation and

specimen necking.

The reduction of area was inversely related to the amount of cold work.

The reduction of area of ductile materials has been associated with fracture

toughness, since the induced triaxial state after necking simulates blunt

notch conditions. The reduction of area for alloys 104 and 155 was inversely

related to yield strength (Fig. 31), like the relationship between fracture

toughness [Kj c (J)] and tensile yield strength (ay) in austenitic stainless

steels at low temperatures, and the slopes of these curves are similar. All

austenitic alloys that failed by ductile rupture in a fracture toughness test

fell within a rather narrow band on a plot of Kj c (J) versus ay, but the

absolute values of the inverse linear relationship for the copper alloys

appear to be distinct. The lack of overlapping lines for alloys 104 and 155

(Fig. 31) implies that either the fracture mode was different (unlikely,

since both alloys fail after much ductility) or that the reduction of area depended on alloying.

Alloy 175 was considerably stronger at all test temperatures than

alloys 155 and 104 in the hard condition. However, the ductility of this

stronger alloy was less, as expected, implying that the toughness of alloy

175 should also be lower than those of alloys 104 and 155.

The yield strength of these copper alloys was mildly dependent on temperature, but strongly dependent on cold work and alloying. Cold- rolling

to the half-hard condition increased the yield strength by about a factor

of 10 for alloy 104 and a factor of 4 for alloy 155. The addition of only

0.06 wt.% phosphorus and 0.11 wt.% magnesium increased the yield strength by

a factor of 2.5. The temperature dependence of the yield strength was also

affected by cold work: Between 76 and 4 K no change in yield strength was measured for the annealed 104 and 155 alloys, whereas the yield strength of

the cold-rolled alloys continued to increase with decreasing temperature.

The work-hardening rate of alloys 104 and 155 in the annealed condition

increased with decreasing temperature (Figs. 21 and 24). Cold-rolling re- duced the amount and rate of work hardening in both alloys (Figs. 22, 23,

25 , and 26)

.

For alloys 155—H and 175—H, tensile measurements were made on specimens

of both longitudinal and transverse orientation with respect to the rolling

direction of the plate. Little effect of rolling direction on the tensile

properties is seen in Tables II and IV. Occasional disparities among the

data of the two orientations are believed to be due to variations in the

material

.

Both displacement- and load-control tests were conducted for alloys

104-1/2 H, 104—H , 155-1/2 H, and 155-H (Figs. 11, 12, 14, 15, and 16).

There was some disagreement in the results for alloys 155-1/2 H and 155-H,

88
but Che tensile properties of the other two alloy conditions were remarkably

consistent; there were no differences due to test control method. Therefore, the disagreement in the data of alloys 155—1/2 H and 155—H is thought

to be due to material variations rather than method of test control.

£1ectrical Resistivity

The electrical resistivities (p) of all conditions of the three copper

alloys were measured at 295, 76, and 4 K and are given in Fig. 32 and

Table V. Table V also includes the residual resistivity ratio (RRR) , which

is the ratio of resistivity at 295 K to that at 4 K. Many consider the RRR

to be a more reliable, relative estimate of material imperfection because it

is independent of specimen area and other absolute measurement inconsistencies. Cold-rolling increased the resistivity of alloys 104 and 155. As a result of annealing, phosphorus and magnesium in alloy 155 were no longer

precipitates, but were solutes, thus increasing the resistivity of the

alloy. Higher strength alloys generally have higher resistivities. The

ratio Oy/p for each alloy and condition is listed in Table V and plotted

versus temperature in Fig. 33. For high- field magnets, designers tend to

select conductors that have high ratios. Alloy 104—H had the best Oy/

p

value at 4 and 76 K, only slightly less than that of alloys 104—1/2 H and

155—H at 295 K. No practical gains are achieved by selecting alloy 175,

unless a very high yield strength is required for the alloy.

The changes in resistivity, Ap, due to changes in condition are shown

in Fig. 34 for alloy 104. Cold-rolling obviously increased Ap. The temperature dependence of the H condition followed that found in a review of

the literature by Simon and Reed;^ that is, Ap changes are larger at higher

temperatures. However, the 76-K data did not follow this trend, and this

disparity likely represents data scatter.

Elastic Properties

The elastic moduli of alloys 104 and 155 were measured at 295, 76,

and 4 K for all conditions. These are discussed in other papers of this

report .

Fatigue and Creep

The creep data for all alloys and conditions are summarized in Fig. 35.

The creep rates are mainly logarithmic, with specimen strain (e) proportional to log t, where t is time. Except for alloy 104—H, the alloys and

conditions have a nearly constant e/log t value. For the annealed alloys

(104 and 155), the slope of e versus log t was lower. Therefore, except for

that of 104—H, the creep rate of the alloys was logarithmic at the stress

level of 0.9tfy.

The fatigue data are summarized in Fig. 36. All alloys have two

distinct ranges of specimen strain versus the log of the number of cycles

(log N) . The early range, at about 10 cycles or less, contains a creep

component and, therefore, exhibits a higher rate of strain. Subsequently,

89
the strain contains only the fatigue -induced component. In both stages,

e is proportional to log N.

A combination of fatigue and creep was also measured for these alloys

.

These data are summarized in Fig. 37. Two stages are apparent when strain

is plotted versus log N. The early stage has a steeper slope (higher e/N)

and is dominated by primary creep strain. For the annealed alloys, the

first stage has a larger slope and extends to larger values of N.

All data are summarized in Fig. 38; the strain components are separated

into initial and total contributions. The contribution of elastic strain is

shown by the solid lines in the figure. All strain contributions scale with

the applied stress. The total strain increased as the applied stress increased. The relative contributions of elastic and plastic strain are shown

more clearly in Fig. 39. The elastic strain (e e ) naturally increased with

applied stress (<7 a ) , following e e = cr a

/E, where E is Young's modulus. The

total plastic strain decreased with applied stress, following the inverse

linear dependence depicted by the solid line in Fig. 39. These dependences

of plastic strain on applied stress were roughly independent of the type of

test: creep, fatigue, or fatigue plus creep. For most tests, alloys 104—

1/2 H and 155—1/2 H had much lower plastic strains than the other alloys

(see data at aa — 250 MPa in Fig. 39). Perhaps these lower strain values

are associated with room- temperature aging effects in cold-rolled copper,

since these two alloys were the "as -received" material and had been rolled

to the half -hard condition about seven years earlier. We are now exploring

ambient temperature aging effects following cold-rolling.

SUMMARY

Tensile, creep, fatigue, combined creep and fatigue, and electrical

resistivity properties at 4, 76, and 295 K were measured for three alloys in

various conditions—annealed, half -hard, and hard. The data were summarized

and compared to determine trends as a function of temperature, alloy condition, time (creep), number of cycles (fatigue), and applied stress (creep,

fatigue, and combined creep and fatigue).

Yield strengths were increased by cold-rolling, but at the expense of

electrical conductivity, ductility, and toughness. Cold-rolling considerably reduced the work-hardening capacity of copper alloys, and it increased

the yield strength much more than the ultimate strength. Therefore, the

increase in yield strength due to cold-rolling is partially negated in

applications where ultimate strength is also important. Furthermore, the

reduced toughness in cold-rolled alloys, combined with the expected higher

operating stresses, places a premium on the alleviation of stress concen- trations in structural design. Since heavily cold-rolled copper has less

strain accommodation than other copper alloys, safety factors must be care- fully evaluated before using these alloys in applications requiring high

yield strength.

On the basis of our tests, alloy CDA 104 in the hard condition had the

best combination of yield strength and electrical conductivity.

90
ACKNOWLEDGMENTS

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

The authors thank Lonnie Scull for cold-rolling the half-hard plate and

Patrick Purtscher for assistance in annealing the half -hard plate. Edward

Dalder of Lawrence Livermore National Laboratory (LLNL) and Charles Bushnell

of Princeton Plasma Physics Laboratory (PPPL) helped us to obtain the copper

alloys. Herbert Becker of the National Magnet Laboratory at MIT assisted in

material selection. Colorado School of Mines kindly allowed us to use their

cold- rolling facilities. The CDA 104 plate was provided by LLNL, who obtained the plate from PPPL. The CDA 155 plate, produced by the Copper Range

Company for fabrication of the B-coil center post for Doublet III, was supplied by General Atomic. The CDA 17510 plate, developed jointly by LLNL,

PPPL, and Cabot Berylco Corporation, was provided by LLNL.

REFERENCES

1. Metals Handbook . 9th edition, vol. 9. Metals Park, Ohio: American

Society for Metals; 1985, p. 401.

2. F. R. Fickett, Electrical properties, chapter 5 in Materials at Low

Temperatures . R. P. Reed and A. F. Clark, eds . Metals Park, Ohio:

American Society for Metals; 1983, p. 166.

3. N. J. Simon and R. P. Reed, Cryogenic properties of copper and copper

alloys, to be published.

4. H. M. Ledbetter and S. A. Kim, Low- temperature elastic constants of

deformed polycrystalline copper, in Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures—

. NBSIR 87-3067,

National Bureau of Standards, Boulder, Colorado; 1987, p. 139.

5. H. M. Ledbetter, Acoustoelastic residual stress measurements: role of

anisotropic dislocation arrays, in Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures—

. NBSIR 87-3067, National

Bureau of Standards, Boulder, Colorado; 1987, p. 151.

6. H. M. Ledbetter and M. W. Austin, Deformed-polycrystalline-copper

elastic constants, in Materials for Magnetic Fusion Energy Applications

at Low Temperatures—

. NBSIR 87-3067, National Bureau of Standards,

Boulder, Colorado; 1987, p. 127.

91
Table I. Copper Alloys

CDA Type Condition Composition (wt.%)

104 Annealed 99.5 [Cu + Ag]

1/2 hard

Hard

0.027 Ag

155 Annealed 99.75 [Cu + Ag]

1/2 hard 0.027 Ag, 0.06 P,

0.11 Mg

17510 Hard 97.6 [Cu + Ag]

1.8 Ni, 0.38 Be

-Table II. Tensile Properties of CDA 175 Copper

TEST

CONTROL

CONDITIONORIENTATION

TEMPERATURE

(K)

TENSILE PROPERTIES

ay

(MPa)

CTu

(MPa)

EL

(%)

R. A.

(%)

DISPLACEMENT H long. 295 715 783 10 24

(e = 1 x 10' 4s’

1

) 76 783 888 15 49

4 814 972 17 35

DISPLACEMENT H transv. 295 718 111 9 34

(e = 1 x 10' 4 s' 1 ) 76 799 890 14 47

4 793 937 18 46

92
Table III. Tensile Properties of CDA 104 Copper

TEST

CONTROL

CONDITION

ORIENTATION

TEMPERATURE

(K)

TENSILE PROPERTIES

°y

(MPa) (MPa)

EL

(%)

R. A.

(%)

DISPLACEMENT A transv. 295 23.0 205 48 87

(e - 1 x 10' 4 s' 1 ) 76 30.5 336 57 87

4 30.5 410 76 82

DISPLACEMENT 1/2 H transv. 295 275 277 13 78

(c - 1 x 10' 4 s' 1 ) 76 315 375 37 82

4 331 446 — —

LOAD 1/2 H transv. 4 326 440 49 81

(P - 1 MPa s' 1 )

DISPLACEMENT H transv. 295 342 357 9 72

(c - 1 x 10' 4 s' 1 ) 76 401 442 19 78

4 431 509 32 72

LOAD H transv. 295 342 356 10 76

(P - 1 MPa s' 1 ) 76 408 452 20 76

4 432 503 30 75

93
Table IV. Tensile Properties of CDA 155 Copper

TEST

CONTROL

CONDI

TION

ORIENTATION

TEMPERATURE

(K)

TENSILE PROPERTIES

ay

(MPa) (MPa)

EL

(%)

R. A.

(%)

DISPLACEMENT A transv. 295 65 244 44 79

(e = 1 x 10’ 4 s’ 1 ) 76 81 378 55 72

4 81 457 62 69

DISPLACEMENT 1/2 H transv. 295 277 306 15 63

(e = 1 x 10’ 4 s’ 1

) 76 307 411 27 53

4 314 479 37 56

LOAD 1/2 H transv. 295 275 310 18 62

(P = 1 MPa s' 1

) 76 300 415 32 63

4 315 479 37 58

DISPLACEMENT H long. 295 383 394 8 58

( e = 1 x 10’ 4

s* 1

) 76 466 520 16 41

4 470 579 32 57

DISPLACEMENT H transv. 295 397 412 8 51

(e = 1 x 10’ 4 s’ 1 ) 76 460 500 19 53

4 494 560 28 57

LOAD H long. 295 360 379 10 61

(P = 1 MPa s’ 1 ) 76 460 524 22 52

4 468 515 15 37

LOAD H transv. 295 388 404 8 51

(P = 1 MPa s’ 1

) 76 460 498 20 52

4 484 548 16 49

LOAD H long. 295 380 403 10 60

(P = 70 MPa s’ 1

) Load H transv. 76 451 496 15 52

(P = 80 MPa s’ 1 )

94
Table V. Electrical Resistivity Data

Alloy T

(K) (ML)

P

(/zQ • cm)

Oy/P

(MPa//jQ cm)

RRR

155-H 4 483 0.2005 2409

76 460 0.398 1156

295 382 1.92 199 9.6

155-1/2 H 4 315 0.177 1780

76 306 0.384 797

295 280 1.955 143 11.0

1 55 - A 4 81 0.358 226

76 81 0.562 144

295 65 2.09 31 5.8

104-H 4 432 0.0392 11020

76 406 0.241 1685

295 343 1.785 192 45.5

104-1/2 H 4 328 0.0278 9892

76 315 0.2195 1435

295 275 1.735 191 61.7

104 - A 4 30.5 0.010 3000

76 30.5 0.206 148

295 23 1.75 0.6 175.0

175-H 4 803 1.50 535

76 791 1.80 439

295 716 3.98 180 29.3

95
Fig. 1. Alloy CDA 104. Form: 32-mm-thick plate.

Condition: annealed at 510°C, 1 h, Ar atmosphere, air cooled.

On face normal to S direction, hardness = 53 Rp; grain size = 80 jim.

96
Rolling direction >

Fig. 2. Alloy CDA 104. Form: 32-mm-thick plate.

Condition: half-hard.

On face normal to S direction, hardness = 80 Rp; grain size = 77 /zm.

97
Fig. 3. Alloy CDA 104. Form: 15-mm-thick plate.

Condition: hard (cold-rolled 50% from half-hard condition).

On face normal to S direction, hardness = 88 Rp; grain size was not

measured owing to extensive texture.

98
T

Fig. 4. Alloy CDA 155. Form: 39-mm- thick plate.

Condition: annealed at 550°C, 1.25 h, Ar

On face normal to S direction, hardness = atmosphere, air cooled.

63 RF ; grain size = 75 /j,m.

99
Fig. 5. Alloy CDA 155. Form: 39-mm-thick plate.

Condition: half-hard.

On face normal to S direction, hardness = 88 R ; grain size = 72 am.

r

100
Fig. 6. Alloy CDA 155. Form: 15-mm-thick plate.

Condition: hard (cold-rolled 60% from half-hard condition).

On face normal to S direction, hardness = 96 R^,; grain size was not

measured owing to extensive texture.

101
T

Fig. 7. Alloy CDA 175. Form: 2

5 -mm- thick plate.

Condition: hard.

On face normal to S direction, hardness = 92 R^; grain size was not

measured owing to extensive texture.

102
Stainless Steel

Fig. 8. Apparatus for displacement-control tests.

103
Applied

Stress,

TEST PROFILES

Fig. 9. Stress profiles for fatigue, creep, and fatigue-plus -creep tests.

104
Strength,

MPa

Fig. 10. The tensile properties of alloy CDA 104—A.

Percent
Fig. 11. The tensile properties of alloy CDA 104-1/2 H.

106

Percent
Fig. 12. The tensile properties of alloy CDA 104—H.

107

Percent
Strength,

MPa

CDA 155-A

0 100 200 300

Temperature, K

Fig. 13. The tensile properties of alloy CDA 155-A.

108

Percent
Fig. 14. The tensile properties of alloy CDA 155—1/2 H.

109

Percent
600

« 500

Q_

£

B)

c

0

CO 400

Temperature, K

Fig. 15. The tensile properties of alloy CDA 155—H.

110

Percent
Strength,

MPa

Fig. 16. The tensile properties of alloy CDA 175—H.

Ill

Percent
1000

1 00 200

Temperature, K

• 175-H

155-H

104-H

155-V2H

1 64-1/2 H

300

Fig. 17. Tensile yield strength summary.

112
Fig. 18. Tensile ultimate strength summary.

113
Reduction

of

Area,

%

'

Elongation,

%

Fig. 19. Tensile elongation summary.

Fig. 20. Tensile reduction of area summary.

114
Stress,

MPa

Stress,

MPa

1000

800

700

600

500

CDA 104 ANNEALED

Fig. 21. Stress—strain curves for alloy CDA 104—A;

displacement-control test.

Fig. 22. Stress—strain curves for alloy CDA 104—1/2 H;

displacement-control test.

115
Fig. 23. Stress—strain curves for ailoy CDA 104—H;

displacement-control test; transverse orientation.

Fig. 24. Stress—strain curves for alloy CDA 155—A;

displacement-control test.

116
Stress,

MPa

700

Fig. 25. Stress—strain curves for alloy CDA 155—1/2 H;

displacement-control test; longitudinal orientation.

Fig. 26. Stress-strain curves for alloy CDA 155-H;

displacement-control test; transverse orientation.

117
Fig. 27. Stress—strain curves for alloy CDA 155—H;

displacement-control test; longitudinal orientation.

Fig. 28. Stress—strain curves for alloy CDA 155—H;

load-control test; transverse orientation.

118
Fig. 29. Stress—strain curves for CDA 155—H;

load-control test; longitudinal orientation.

Fig. 30. Stress-strain curves for CDA 175-H;

displacement-control test; transverse orientation.

119
Fig. 31. Reduction of area as a function of yield strength

for alloys 104 and 155.

120
Resistivity,

cm

Fig. 32. Resistivity as a function of temperature.

121
10000

1000

100

10

Fig. 33. Ratio of yield strength to resistivity as a function of temperature.

Condition

Fig. 34. Changes in resistivity due to changes in alloy condition.

122
Fig. 35. Summary of creep data at 295 K; a = 0.9 a^.

Fig. 36. Summary of fatigue data at 295 K; o - 0.9 a .

123
6000 -

5000

10° 10 1

175-H

102

Cycles

io3 io4

Fig. 37. Summary of combined creep and fatigue data

at 295 K; o = 0.9 a . y

124
6000

5000

4000

3000

2000

1000

0

6000

5000

4000

3000

2000

1000

0

6000

5000

4000

3000

2000

1000

elastic strain

o - total strain

• - initial strain

. Fatigue

elastic strain

Applied Stress, MPa

Fig. 38. Summary of all creep and fatigue data at 295 K.

125
Fig. 39. Relative contributions of elastic and plastic strain versus stress.

126


DEFORMED -POLYCRYSTALLINE -COPPER ELASTIC CONSTANTS*

H. M. Ledbetter and M. W. Austin

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

By ultrasonic methods, we studied the principal macroscopicorthotropic- symmetry elastic constants of plastically deformed

copper plate. By x-ray diffraction, we verified a strong plate

texture, which we identified as (110) [112, ill]. From the monocrystal elastic constants, we calculated the macroscopic elastic

constants expected for ideal textures. Prediction failed to fit

observation; mostly, predicted elastic stiffness was too high.

We ascribe this disagreement to a high density of deformationinduced dislocations. To the elastic-constant measurements,

these dislocations contribute a reversible plastic strain. For

the elastic shear stiffnesses, the usual macroscopic

symmetry failed. We ascribe this breakdown to dislocation-array

anisotropy. For the elastic constants, the dislocation contribution overwhelms the texture contribution. After plastic deformation, the elastic constants show an ambient- temperature time

dependence

INTRODUCTION

In 1974, Ledbetter and Naimon [1] reviewed copper's elastic constants,

both polycrystal and monocrystal. They included effects of temperature and

pressure. Plastic-deformation effects they treated only cursorily for two

reasons: complexity and dearth of careful systematic studies. Texture

effects they omitted.

At least since Gruneisen in 1907 [2]we know that plastic deformation

changes apparent elastic moduli, sometimes considerably. Thirty- four years

later, Lawson [3] connected the Young-modulus decrease in compressed copper

with Taylor dislocations. Models relating elastic-stiffness decrease to

dislocation vibrations appeared from several sources: Eshelby [4],

Koehler [5], Friedel [6], and Granato and Lucke [7].

Intended for publication in Phvsica Status Solidi .

127
We remain uncertain about who first studied texture effects and elastic

constants. Probably this occurred a century ago. In 1889, Voigt [8] considered the effective elastic constants of a random polycrystalline aggregate.

Probably, nonrandom (textured) aggregates were considered then.

In recent years, Ledbetter [9,10] reported the elastic constants of

annealed quasi- isotropic polycrystalline copper. Using various models,

he related polycrystal and monocrystal elastic constants and found that

Kroner's model [11] works best.

The present study considers the macroscopic elastic constants of

polycrystalline copper deformed by rolling. We consider effects of both

texture and plastic deformation.

We studied two commercial copper alloys: CDA104, oxygen- free copper

with silver; and CDA155, silver-bearing copper. Including silver, these

alloys contain 99.95 and 99.75 percent copper, respectively. We studied

three plastic - deformation conditions: the received deformed plate (hard- ness =80, 88 Rp) , plate deformed by rolling to a 60 percent reduction

(hardness =88, 96 Rp) , and plate in the recrystallized-annealed condition

(hardness =53, 63 Rp) . From these, we ground parallelepiped specimens

measuring approximately 1.5 cm for all three dimensions. Figure 1 shows

a typical microstructure. We obtained the annealed condition by heating the

received plates for 1 h at 510°C and for 1.3 h at 550°C, respectively.

MEASUREMENT METHODS

Mass density, p, we measured by Archimedes's method using distilled

water as a standard. We estimated inaccuracy as 0.05 percent.

Sound velocity, v, we measured by a MHz - frequency pulse-echosuperposition method described previously [9,10,12]. From mass density

and sound velocities, we computed elastic constants by

Here, subscripts S. and t denote longitudinal and transverse. For each of

six materials, we measured six sound velocities corresponding to the longi- tudinal elastic constants C 11( C 2 2 » and C 33 and the transverse (shear)

elastic constants C 44 , C 5S , and C 66 . Because they contribute little to

the present study and because they contain larger errors , we omitted

measurements of C 12 , C 13 , and C 23 . For sound velocities, we estimated a 0.2 percent inaccuracy in v^ and a 0.4 percent inaccuracy in v^. Thus, we estimated errors as 0.45 percent and C t errors as 0.85 percent.

MATERIALS

( 1 ) and

( 2 )

128
Texture we studied by x-ray diffraction from surfaces perpendicular

to the three principal directions: x x = rolling direction, x3 = direction normal to rolling plane, x2 = direction orthogonal to x x and x3 (some

call x 2 the long- transverse direction). For x-ray diffraction, we used a computer-controlled commercial diffractometer and Cu-Ka radiation.

RESULTS

The study's principal results occur in Table 1. The table shows

measurements for two alloys in three plastic-deformation conditions. For

one alloy, CDA155, in the high-hardness condition, we show measurements at

two post-deformation times: 15 days and 105 days.

Table 1 also shows results calculated from copper's monocrystal elastic constants: C 1X = 168.4, C 12 = 121.4, and C 44 = 75.39 GPa. These

calculations include the quasi- isotropic elastic constants by Kroner's

method [11] by Voigt's method, elastic constants for three textures:

(110) [ 112 ] , (110) [111] , (112) [111] . The fourth texture, (110) [ 112 , ill],

equals an arithmetic average of the first two.

Figure 2 shows a set of four typical x-ray-diffraction patterns:

standard copper and patterns from the x x , x2 , x3 faces. Study of these

patterns led to the following interpretation. With respect to annealed

copper, diffraction lines are broadened but little shifted. For the x x and x 2 faces, one sees essentially identical patterns. These differ

dramatically from the x 3 -surface pattern. It shows no (111) line and a strong (220) line. Unambiguously, this shows that {110} corresponds to

the rolling plane. We expect a rolling direction of either <11 1> or <112>.

Diffraction patterns from the x x and x2 surfaces suggest a mixture of both,

favoring <1 1 1>

PLATE -TEXTURE ELASTIC CONSTANTS

We denote a plate texture by the notation (hki) [uvw] , where (hki)

denotes the rolling plane and [uvw] the orthogonal rolling direction. Here, we take x x = rolling direction, x2 = transverse direction, x 3 = rollingplane normal. In an ideal plate texture, only [uvw] directions are parallel

to x x and only (hk/0) plane normals are parallel to x3 . Thus, in many ways, an ideal plate texture behaves like a monocrystal. To obtain the plate

texture elastic constants, Djjki> we perform the usual tensor-property co- ordinate transformation:

^ijki = a imajnakoaip^mnop (3)

Here, Cmn0p

denotes the usual fourth- order monocrystal elastic constants and

the a^m

denote the direction cosines between the new (plate) co-ordinates

and the old (crystal) co-ordinates.

129
For (110) [112] plate texture, the new base vectors are x x = [112], x 2 =

[111], and x 3 = [110]. In Voigt's contracted notation, for this case, one obtains the six-by-six matrix (for simplicity, we show only indices):

11 12 13 0 0 16

12 22 12 0 0 0

13 12 11 0 0 -16 . (4)

0 0 0 44 -16 0

0 0 0 -16 (11—13)/2 0

16 0 -16 0 0 44

Here

,

Du = (Cn + Cl 2 + 2C 44 )/2

D 2 2 O

'

' II + 2Ci 2 + 4C 4 4 >/3

Dl 2 - (C 11 + 2C X 2 - 2C 4 4 )/3

Dl 3 = (Cn + 5C X 2 - 2C 4 4 )/6

Dl 6 = (Cl! - Ci 2 - 2C 4 4 )/yi8

D 44 = (Cl! - Ci 2 + c 44 )/3.

For (112) [111] plate texture, x 1 = [111], x 2 one obtains

[liO] , x 3 = [112] , and

22 12 12 0 0 0

12 11 13 0 16 0

12 13 11 0 -16 0

0 0 0 (11—13)/2 0 16

0 16 -16 0 44 0

0 0 0 16 0 44

130
Here, the Djj elements are given by Eq. (5). To compute the D^j elements,

for copper's monocrystal elastic constants we used C X1 = 168.4, C 12 = 121.4,

and C 44 - 75.39 GPa [1]

For these two texture cases—(110) [112] and (112) [111], the results

are given in Table 1. Also, the table shows results for a (110) [112, 111]

texture suggested by our x-ray-diffraction results: a (110) rolling plane

and a rolling direction consisting of an approximately equal mix of [112]

and [111] directions To make this mixed-case prediction, we averaged

(110) [112] case with the (110) [ 111

case where x x = [111], x2 = [112], x 3 - [ 110] , and

22 12 12 0 0 0

12 11 13 0 0 -16

12 13 11 0 0 16

D ij “

0 0 0 (11—13)/2 16 0

0 0 0 16 44 0

0 -16 16 0 0 44

The D-jj elements are given by Eq. (5).

DISCUSSION

Results in Table 1 show several features:

1. For a rolled plate, the C^-j show the expected orthotropic symmetry:

Ci i ^ C 2 2 ^ C 33 and C 44 ^ C S5 ^ C 6 6 .

2. Excepting C 232 3 and C 2121 , all the principal-diagonal elastic stiffnesses are lower than those predicted for (110) [112, 111] texture.

3. For the elastic shear stiffnesses, the usual C ijki =

^kJij symmetry

fails.

4. With time, all the C^-s increase: over three months, 1 percent in the

(i = 1,2,3), and 2 percent in the (i = 4,5,6).

The first feature, macroscopic anisotropy, results from texture; and,

as described below, from anisotropic dislocation arrays.

We explain the unexpected second and third features by a single hypothesis: an anisotropic high-density dislocation array. Dislocations arise

131
naturally during plastic deformation; heavy deformation induces large

numbers. In fee metals, the principal dislocations lie on {111} planes in

<110> directions. Because of texture, these planes and directions orient

preferredly rather than randomly. For the present case, we expect {111}

planes perpendicular to both x x and x2 and <110> directions parallel to x 3

.

Perhaps Lawson [3] first discovered the relationship between dislocations and elastic constants. He applied compressive stress to polycrystalline copper and found a Young-modulus decrease of six to seven percent.

A theoretical basis for dislocation- lowered elastic constants arose from

Eshelby [4] . He showed that the shear-modulus varies according to

SG' 1 G* 1

( 8 )

Here, n denotes dislocation density, a unit-cell dimension; d, dislocationmotion distance, and aQ , elastic shear stress. For a pinned dislocation

under oscillating stress, Koehler [5] gave a detailed calculation. He

showed that reversible plastic dislocation motion creates a small extra

strain that lowers the Young modulus. In measurements on copper monocrystals, Thompson and Holmes [13] found that mechanical deformation reduces the

Young modulus up to ten percent. Calculations by Koehler and deWit [14]

produced three conclusions: (1) elastic anisotropy increases the dislocation-network elastic-constant lowering; (2) for annealed copper, the effect

equals a few percent; (3) for slightly deformed copper, the effect may reach

ten percent. Extending Koehler's vibrating- string model, Granato and Lucke

[7] gave a detailed dislocation-damping model that contains an associated

modulus change, AE/E, with each loss mechanism involved in dislocation

motion. These authors showed that AE/E depends on both strain amplitude

and frequency (at high frequencies).

From these five studies, both experimentally and theoretically, we expect unpinned dislocations to lower the elastic stiffness, C. The

following relationship summarizes the main idea:

C =

e + e e p

(9)

Here, a denotes stress; e, strain; e, elastic; and p, plastic. For wellpinned dislocations, ep equals zero. Pinning results from many mechanisms

that include impurities, irradiation- induced defects, and dislocation nodes.

We now consider whether basic physical relationships allow such a large

modulus decrease. For low strain amplitudes, Koehler [5] derived an expression for the shear-modulus change

dG

G

NGa5 2CL3 c

2 ( 10 )

Here, N denotes atomic lengths of dislocation line; G, shear modulus; a, interatomic spacing; L

3

, specimen volume; c, concentration of pinning points

132
along dislocation line (n/N) ; and 7rC , the line tension that tends to

decrease dislocation length. According to Mott and Nabarro [15]

r 2Ga 2

. (1 - v)

Here, v denotes Poisson ratio. Combining Eqs . (10) and (11) gives

dG _

7rPa 2 (1 — u) . G 4c 2

(11)

( 12 )

We substituted D = Na/L3

, D being the usual dislocation density (total

length per unit volume, cm/cm3 ). From Eq. (12), we can compute the dislocation density required to give a ten-percent shear-modulus decrease. For

copper, we take a = 2.56 A and u = 0.345. We assume c = 10' 2

, one pinning

point per hundred atoms. Then, we obtain a reasonable number:

D = 3- 10 10 cm/cm3

. (13)

According to Hull and Bacon [16], well-annealed metal crystals show a D of

10 6 to 10 8 cm" 2 and heavily deformed metals show D values up to 5-10 13 cm' 2

. We conclude that the observed modulus decrease agrees with existing models.

By using a model of Seeger and Haasen [17], we can estimate the dislocation density in our deformed materials. For nondeformed crystals, these

authors concluded that dislocations increase volume by one atomic volume, Q,

per length of dislocation equal to the Burgers -vector magnitude, b. Thus,

D — b.

p'Q (14)

Here, p denotes mass density, and we found —dp/p = 2.5-10' 4

. Taking b =

2.56 A and = 11.81 A

3

, we obtain D = 5-10 11 cm/cm3

. This agrees within

one order of magnitude with the estimate of Eq. (13). Considering contributions from point defects would give closer agreement.

It remains to consider the dislocation-array anisotropy. In fee crystals such as Cu, Ag, and Au, the principal dislocation is an (a/2)<110>

Burgers -vector edge dislocation that glides on {111} planes and lies along

<11 2> directions [18]. In a nontextured aggregate, all these vectors are distributed randomly, and the 6p term in Eq. (9) is isotropic. Texture

causes a nonrandom distribution of these dislocation vectors and makes Cp

anisotropic. Figure 3 shows the source of this anisotropy by considering a single edge dislocation in an isotropic block. If we assume that the pinned

dislocation can move easily in its glide plane, but not otherwise, then it

follows for the longitudinal elastic moduli that

^1111 = ^3333 > ^2222' ( 15 .)

Similarly, for the shear moduli, it follows that

C'2 323-> C'3232> (16)

133
C 2 1 2 1 > ^1 2 1 2 > and

^1313 = ^ 3131 •

Here, the first index indicates shear-plane normal and the second index

indicates shear direction. For this simple case, Eqs„ (15) through (17)

represent an anisotropic dislocation contribution to the elastic constants

through an anisotropic ep. By considering screw dislocations and partial

dislocations, one finds corresponding anisotropies that break the usual

C^j symmetry for shear elastic constants and the C 1X11 = C 2222 =

C 3

3

3 3 symmetry for longitudinal elastic constants. We expect the largest

symmetry breaking for the shear constants. The shear constant showing the

largest effect corresponds to the shear-wave polarization direction con- taining the Burgers vector, x2 in Fig. 3.

Results in Table 1 for the shear constants show that symmetry breaking

occurs for C 44 and C 66 , but not for C 55 . This requires that the Burgers

vectors favor x2 , which is consistent with the texture result that <110> directions lie along x 2

. The remarkably low C 22 value for the CDA155 alloy

also supports the suggestion that Burgers vectors lie along x 2.

Although we failed to study the phenomenon systematically, we found

that after plastic deformation all the Cji increased with time. Over three

months, the’ increase was 1 to 2 percent, being larger for the shear constants

C 4

4

, C 55 , and C 66 . We interpret this recovery to result from impurity atoms

diffusing to and pinning dislocations, thus reducing £p in Eq . (9).

Finally, we discuss texture effects. As described in the results sec- tion, x-ray-diffraction results from the three faces orthogonal to x x , x 2 , and x 3 , when compared with those of nontextured copper, show a strong texture: a { 110 } rolling plane and a rolling direction that combines <110> and

<112>. From the calculated results in Table 1, we expect this texture to

increase C xx , C 22 , and C 33 , decrease C 66 , and leave C 44 and C 55 unaffected.

Observation fails to confirm these expectations. No other reasonable texture model gives results that agree with observation. We conclude that

texture effects obscure the dislocation effects described above. In studying deformed-copper-rod elastic constants, Wawra [19] reached a similar

conclusion: "texture ... is not decisive in our problem."

(17)

(18)

CONCLUSIONS

From this study, there emerged six conclusions:

1. In rolled copper plate, elastic constants depend on two conditions:

texture and plastic deformation.

2. For heavily deformed materials, plastic-deformation effects can obscure

texture effects. For both longitudinal and transverse elastic constants,

plastic deformation may lower the elastic stiffness by up to 15 percent.

134
3. Elastic stiffne'ss is lowered by a high density of nonpinned dislocations,

that is, a reversible plastic strain.

4. The macroscopic symmetry breaks because of an anisotropic

dislocation array.

5. After plastic deformation, elastic stiffness increases with time because

the density of nonpinned dislocations decreases, caused perhaps by

impurity diffusion to dislocations.

6. Annealing restores elastic constants that typify the nondeformed texturefree quasi- isotropic state.

ACKNOWLEDGMENT

The U.S. Department of Energy, Office of Fusion Energy, provided support.

S. A. Kim made the velocity measurements. L. Scull and R. Walsh furnished

samples and metallography. Professor E. Kroner (Stuttgart) made many

valuable comments. M. Lei, guest worker from Institute for Metal Research

(Shenyang) verified the matrix equations.

REFERENCES

1. H. M. Ledbetter and E. R. Naimon, J. Phys . Chem. Ref. Data 3., 897

(1974).

2. E. Gruneisen, Ann. Phys. 22., 801 (1907).

3. A. W. Lawson, Phys. Rev. 60, 330 (1941).

4. J. D. Eshelby, Proc. Roy. Soc. Lond. 197 . 396 (1949).

5. J. S. Koehler, in Imperfections in Nearly Perfect Crystals, Wiley,

New York (1952) , p. 197.

6. J. Friedel, Philos. Mag. 44, 444 (1953).

7. A. Granato and K. Liicke , J. Appl. Phys. 27., 789 (1956).

8. W. Voigt, Ann. Phys. 38, 573 (1989).

9. H. M. Ledbetter, J. Phys. D: Appl. Phys. 12, 1879 (1980).

10. H. M. Ledbetter, Phys. Status Solidi (a) 66., 477 (1981).

11. E. Kroner, Z. Phys. 151, 504 (1958).

12. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl. Phys. 51 . 305 (1980).

13. D. 0. Thompson and D. K. Holmes, J. Appl. Phys. 27, 191 (1956).

14. J. S. Koehler and G. DeWit, Phys. Rev. 116, 1121 (1959).

15. N. F. Mott and F. R. N. Nabarro, in Report on a Conference on the

Strength of Solids, Physical Society, London (1948), p. 8.

16. D. Hull and D. J. Bacon, Introduction to Dislocations, Pergamon, Oxford

(1984), p. 22.

17. A. Seeger and P. Haasen, Philos. Mag. 3, 470 (1958).

18. A. Seeger, in Dislocations and Mechanical Properties of Crystals,

General Electric, Schenectady (1957), p. 243.

19. H. Wawra, Phys. Status Solidi (a) 54, 291 (1979).

135
Table

1.

Calculated

and

measured

elastic

constants

pf

annealed

and

deformed

copper

alloys,

units

GPa

O' <3

m r- o K* m ON CM CM CM CM o

CM o 00

O >3 in >3* *3 *3 CM 4-H CM

O CM ON

r-*

O

m 00 4—4 >3

00 ^H CO 4-H r—4 T—4O'

»-H o 00 o >3m *3 <3 CO

in i CM 4-H CM m v£> co

On < O'

m 4-H 4-H

m

Q ^3

<r o >3 ^H ^3 r*^ 00

m o o O' in uo *4* <3 >3 CM CM

00

H

CO

i—4 vO uo CM UO

w >3 u »—4

4-H 4-H 00 00 00 oo 00

1—4 o o o <3 ^3 «3 >3 >3 3 CM CM CM

CO co CM

<D M3

o£ 00

>3

Q) i—4 v£> 00 U 00 3 <r m CO <3 w CO O' CM oo ^3 <3 >3 >3

ca

O O' O' *3

0) CM 4—h 4-H

O s 00

CO uo

in 4-H v£> O' CO >3

<r i—4 o r*^ CO v£> VO o o o O' <3 <3 >3 •*3 <3 CM CM < o CO Q 00 O vO

CO *3

O

O >3 >3 CM CM CM 00 oo 00 CO 00 o o o ^3 *3 "3 ^3 <r CM CM CM

CO >3 m

00

>3

i«—(

i*—

m uo oo

i—4 a> ON O' o O' O' o ^ - 3 CM CM CM 3 <3 <3 O CM •u CM CM CM

«-H i 4—

X

4-H 4—

H

a>

w

—

H *

CO 0) oo oo u

•-—( ^ - 3 00

O o 00 o o 3 CM •«-» •u CO CM CM m >3 >3

CO 4—1 *—1 X CM CM CM

<u *H r— <D —' H

<u 0) 4-h CO 00

4->

CCJ ^ —• 3 00 o O co

O o O >4-4 •u CO CM CM m >3 >3 3 4—4 «— X CM CM CM a 4—4 4—4 0)

*--4

w

—

H

CJ

0) 00 i—4 OO

V4 ^ — 3 o CO

O o 00

O

O CM

u CM CO CM uo >3

4—4 44—

X CM CM CM

4-H i— a) W — Ho CM CM CM

»H

CL 4-H 4—4 i-4 00 oo ao O o o o <r >3 <3 u CM CM CM uo

CO

4-4

CO CM CO i—4 CO CM

CO CO CM CO i—4 T—4

CM

O a) CM CO 4-4 CO CM i—4

T-l

C CO CM 00 r—« « - 4 CM

u

u 4-1 '-Z v,_v 00 CO 4-1 iH 03 H-l —1 04 00 to « O —1 iJ -4 N 00 w0 CD > u CO o CJ o O o U

o>co

uo

r*.

*3IM

CJ

CO

•-HI oco

•o

0)

cn

<3 cQ K

136

First

number

Indicates

Rp

hardness.

Second

number

indicates

different

times

after

plastic

deformation.

The

hardnesses

represent

three

mechanical

-deformation

states:

annealed,

half-hard,

and

hard.
250 pm

Fig. 1. Microstructure of CDA104 alloy in Rp = 88 hardness condition.

>-

COz

LUH

Z

JL

o o o o CM o CNJ £ l5_

40 65 90

X3,X2 —f x3’ x 1

^x3

CM

CM

CM 115

x1> x 3

co ^

Standard

140

20

Fig. 2. X-ray-diffraction patterns for CDA104 alloy in Rp = 88 hardness

condition. Top two patterns represent x 3 face, rotated around x 2 and rotated around x x

. Third and fourth patterns represent x 2 and x x faces, both rotated around x3

. At the bottom, we show a standard pattern for annealed isotropic copper.

137
Fig. 3. An ideal block containing a single edge dislocation and showing

propagation directions and polarization directions for various

elastic shear waves corresponding to C 232 3 and C 3232 . For the

six principal elastic constants, the dislocation breaks usual

macroscopic symmetries, causing C 111X = C 3333 > C 2222 and C 1221 = ^1313 = ^1331 = C 3223 > C 3232 - C 1212 .

138


LOW -TEMPERATURE ELASTIC CONSTANTS OF DEFORMED POLYCRYSTALLINE COPPER*

H. M. Ledbetter and S. A. Kim

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

By ultrasonic methods, we studied copper's elastic constants

between 295 and 4 K. The specimens consisted of polycrystals

deformed by rolling. At ambient temperature, compared with an- nealed texture-free copper, the deformed specimens usually showed

lower elastic stiffness: up to 11 percent in Young modulus and

19 percent in shear modulus. The Poisson ratio increased up to

14 percent. These softenings reflect three internal - structure

changes: texture, nonpinned dislocation density, and anisotropic

dislocation array. Cooling to 4 K showed regular behavior:

nearly the same as that of annealed texture-free copper. After

plastic deformation, at ambient temperatures, the elastic con- stants showed time -dependent recovery.

INTRODUCTION

Elastic constants vary measurably with common metallurgical variables:

temperature, stress (or pressure), and alloying. Unless a phase transforma

tion intervenes, a material's physical properties usually vary smoothly and

monotonically . Often, one can successfully model physical-property changes

associated with these variables.

Plastic deformation presents another case. It changes the internal

structure: texture, dislocation density, dislocation arrangement, pointdefect density, internal strain, and so on. These internal variables may

vary nonmonotonically with plastic deformation. They may depend on secondary factors: deformation mode (tension, compression, rolling, and so on),

deformation rate, and time after deformation (recovery). Unlike the variables listed above, microscopic plastic deformation is difficult to specify

Usually, we resort to macroscopic average specifications: elongation,

reduction- in- area, hardness, and so on.

* Intended for Materials Science and Engineering .

139
Experimentally, plastic-deformation effects on elastic constants are well-known. Lawson [1] applied compressive stress to polycrystalline copper

and found a Young-modulus decrease of 7 percent. On copper monocrystals,

Thompson and Holmes [2] found Young-modulus decreases up to 10 percent.

In copper deformed at 78 K, van den Beukel and Brouwer [3] found a shearmodulus decrease of 18 percent. There have been numerous similar studies,

too many to describe (for examples, see references in reference 4).

Many theoretical models based on dislocations predict that the apparent elastic stiffness decreases as the density of nonpinned dislocations

increases. Sources of these models include Eshelby [5], Koehler [6],

Friedel [7], and Granato and Liicke [8]. Calculations by Koehler and

deWit [9] produced three conclusions: (1) elastic anisotropy increases

the dislocation- induced elastic-constant lowering; (2) for annealed copper,

the effect equals a few percent; (3) for slightly deformed copper, the

effect ranges up to ten percent.

From the above experimental and theoretical studies, we expect non- pinned dislocations to lower elastic stiffness, C. The following relationship summarizes the main idea:

e (e) + e (p)

( 1 )

Here, a denotes stress; e, strain; e, elastic; and p, plastic. The plastic

strain is reversible, resulting from dislocations bowing between pinning

points. For well -pinned dislocations, equals zero. Pinning results

from many mechanisms that include impurities, irradiation- induced defects,

and dislocation nodes.

This study considers the effects of ambient- temperature plastic

deformation (rolling) on the elastic constants of polycrystalline copper.

By ultrasonic methods, we measured the six principal C-n (C^p, i = 1,6)

between 295 and 4 K. We report values for the three principal Young moduli

(E x , E 2 , E 3 ) and the principal shear moduli (G 23 , G 13 , G 12 ).

MEASUREMENTS

Materials

We studied two commercial copper alloys: CDA104 and CDA155. Including

silver, these alloys contain 99.95 and 99.75 percent copper, respectively.

For each alloy, we studied three plastic - deformation conditions: the re- ceived deformed plate (hardness =80, 88 RF ) , a plate deformed by rolling

to a 60 percent reduction (hardness =88, 96 Rp)

,

and the recrystallizedannealed condition (hardness =53, 63 Rp)

.

From these, by grinding, we prepared parallelepiped specimens measuring 1.5 to 2 cm. Figure 1 shows

a typical microstructure. We obtained the annealed condition by heating

the received plates for 1 h at 510°C and for 1.3 h at 550°C, respectively.

140
Ambient-Temperature Methods

Mass density, p, we measured by Archimedes's method using distilled

water as a standard. We estimate inaccuracy as 0.05 percent.

Sound velocity, v, we measured by a MHz - frequency pulse-echosuperposition method described previously [10]. From mass density and sound

velocities, we computed elastic constants by

C 2 ” (2a)

and

C t - pv£. (2b)

Here, subscripts H and t denote longitudinal and transverse. For each of

six materials (two alloys, three hardnesses), we measured six sound velocities corresponding to the longitudinal elastic constants C 11 , C 22 , and

C 33 and the transverse (shear) elastic constants C 44 , C S5 , and C 66 . Because they contribute little to the present study and because they contain

larger errors, we omitted measurements of C 12 > C 13 , an<i C 23 . F° r sound

velocities, we estimated a 0.2 percent inaccuracy in v^ and a 0.4 percent

inaccuracy in vt

. Thus, we estimated errors as 0.45 percent and C t errors as 0-. 85 percent.

Low-Temperature Methods

We achieved measurements at liquid-helium temperature by methods

described previously [11]. At each temperature, at approximately 5-K

intervals, we made absolute Vj? and two vt measurements in three coolings

of the same specimen. The two vt measurements correspond to shear-wave

polarization along the two principal directions.

Texture

Texture we studied by x-ray diffraction from surfaces perpendicular

to the three principal directions: x 3 = rolling direction, x 3 = direction normal to rolling plane, x2 = direction orthogonal to x x and x3 (some

call x 2 the long- transverse direction). For x-ray diffraction, we used a computer-controlled commercial diffractometer and Cu-Ka radiation.

RESULTS

The study's principal results occur in Table 1. The table shows

results for two alloys in three plastic-deformation conditions. For one

alloy, CDA155, for the high-hardness condition, we show measurements at

two post-deformation times: 15 days and 105 days. For comparison, Table 1 shows results for annealed texture-free copper. For each material, we mea- sured the principal C^j (C-j^, i = 1—6). To compute the Young moduli, E^,

and the Poisson ratios i/jj , we inverted the six-by-six C-jj matrix to obtain

the S^j matrix. This requires knowing the three off-diagonal C^j

which we

141
did not measure. We estimated them under the assumptions that neither

texture nor dislocations affect the bulk modulus

B = (C u + 2C 12 )/3. (3)

Then, one can compute the Young moduli by

Ei = 1/Sii (M

and the Poisson ratios by

"ij - “5^- (5)

IX

Figure 2 shows the temperature variation of two of the elastic con- stants, a longitudinal and a shear. We fit these results to Varshni's [12]

relationship

:

C(T) = C(0) - s/ [ exp ( t/T) - 1]. (6)

Equation (6) contains three adjustable parameters: C(0) , the elastic stiff- ness at zero temperature; s, which relates to zero-point energy; and t, the

Einstein temperature. Compared with reference copper, we found that the

deformed coppers showed essentially identical C(T) behavior when one superimposes the curves.

For T = 4 K, Table 2 nearly replicates Table 1. Table 3 gives the

Varshni-fit parameters, together with those reported previously for oxygenfree high- conductivity copper [13].

Figure 3 shows a typical set of four x- ray- diffraction patterns: standard copper and patterns from the x : , x 2 , x 3 faces. Study of these patterns

led to the following interpretation. Diffraction lines are broadened but

little shifted. From the x x and x2 faces, one sees essentially identical

patterns. These differ dramatically from the x 3 -surface pattern. It shows

no (111) line and a strong (220) line. Unambiguously, this shows that {110)

corresponds to the rolling plane. We expect a rolling direction of either

<1 1 1> or <112>. Diffraction patterns from the x x and x 2 surfaces suggest

a mixture of both, favoring <111>.

DISCUSSION

To begin, we consider the results in Table 1. For the Young moduli,

Ep, and the shear moduli, G^j

,

almost always, the deformed elastic stiffness

is less than the annealed-recrystallized value (hardnesses 53 and 63)

.

For

the Young modulus, the elastic softening ranges up to 11 percent; for the

shear modulus up to 19 percent. This softening results from a high density

of nonpinned dislocations. These dislocations contribute a reversible plastic component to the strain.

142
Unlike the quasi- isotropic cases where Ej_ and G^j are independent of

direction, in the deformed material, E x * E 2 * E 3 , and G 2 3 ^ G 13 ^ G 12 . This emphasizes the orthotopic macroscopic symmetry of the deformed material: a rolled parallelepiped.

Although we failed to study time effects systematically, we found that

at ambient temperatures, the softening tends to disappear and elastic stiffness tends to approach the reference values. Several studies [4] found such

recovery, usually ascribed to dislocation pinning by point defects. For our

few measurements, we observe that the dilatational elastic constants recover

more than the shear elastic constants.

Concerning texture, which we verified by x-ray diffraction, we gave a thorough analysis elsewhere [14], The effect of (110) [ 1 1 2 , 111] texture is

to increase E x , E 2 , G 23 , and G 13 . It decreases E 3 and G 12 . Because these

effects disagree with observation, we concluded that texture affected anisotropic elastic stiffness only secondarily.

The most remarkable result shown in Table 1 is the breaking of the

usual symmetry

^ijkl =^ij lk- ( 7 )

We see that G 2323 ^ G 2332 , G 1313 G 1331 , and G 1212 ^ G 1221 . With a measurement uncertainty of less than 1 percent, we found differences up to

17 percent that averaged 5 percent. This symmetry breaking means a nonsymmetry of in the usual expression for Hooke's law:

CT ij

~= c ijkl e kl • ( 8 >

In Eq. (1), we wrote that the total strain consists of two parts:

e kl + € (P)

kl (9)

This nonsymmetry arises naturally if we consider a single edge dislocation

(Fig. 4) lying along x x with Burgers vector, b, along x2

. When shear waves

pass the dislocation, it can bow between pinning points, bow in the glide

plane, but not out of the plane. If a shear wave passes that propagates

along x 3 and is polarized along x2 , the dislocation contributes to

If a shear wave passes that propagates along x 2 and is polarized along x 3 , the dislocation contributes nothing to e 23 . Dislocations contribute zero

strain when b’g = 0, where g denotes the polarization vector. In terms of

elastic constants, this means C 2323 > C 2332 . In the notation in Table 1,

G 2 3 ^ G 3 2 .

Thus, we understand the shear-modulus symmetry breaking when the

deformed polycrystalline materials contain anisotropic dislocation arrays:

arrays where b fails to average to zero over direction. Elsewhere [15], we gave further analysis on this point. There, we emphasized the importance

of the present results for acoustoelastic residual-stress measurements.

143
Measuring residual stress typically requires detecting a shear-wave velocity

difference of 1 in 10 5

. Obviously, anisotropic dislocation arrays caused

by plastic deformation could obscure and confuse shear-wave-birefringence

residual- stress measurements.

Interestingly, in all the deformed materials the Poisson ratio increased, up to 13 percent. For an isotropic material, the Poisson ratio u relates simply to shear modulus G and bulk modulus B:

1 3B ~ 2G

2 3B + G

( 10 )

By taking v ~ 1/3 and differentiating, one obtains

_ 4 fdB _

dG'

i/ 9 [ B G^

Thus, accompanying a large decrease in G and a small change in B, we expect

a large increase in u . Usually, changes in v reflect changes in interatomic

bonding [16]. In the present study, i/ increases reflect increased nonpinned

dislocation density.

CONCLUSIONS

From this study, there emerge five principal conclusions:

1. In copper deformed heavily by rolling, elastic constants softened

considerably: up to 11 percent in the Young modulus. The softening resulted principally from an increased density of nonpinned

dislocations

.

2. The (110) [ 1 1 2 , 111] texture, which increases six of the nine principal

Young and shear moduli, played a secondary role.

3. Shear-modulus symmetry breaking, for example G 232 3 * G 2332 , occurred

for all the shear moduli. We hypothesize that this asymmetry arises

from an anisotropic dislocation array. This array contributed a reversible plastic strain that varied with deformation mode.

4. At ambient temperature, the elastic stiffnesses recovered toward their

annealed, texture -free values. We ascribe this time dependence to

dislocation pinning by point defects.

5. The low- temperature results presented no surprises. The various E^(T)

and G^j (T) curves almost superimpose on those for nondeformed copper.

144
ACKNOWLEDGMENT

For this study, the DoE Office of Fusion Energy provided financial

support. Using materials provided by Lawrence Livermore National Laboratory

and Princeton Plasma Physics Laboratory, L. Scull and R. Walsh provided

material preparation. M. Austin provided elastic-constant-measurement

guidance

REFERENCES

1. A. W. Lawson, Phys . Rev., 60 (1941) 330.

2. D. 0. Thompson and D. K. Holmes, J. Appl. Phys'., 27. (1956) 191.

3. A. van den Beukel and C. Brouwer, Philos. Mag., 17. (1968) 453.

4. G. Fantozzi, D. Boulanger, and P. Gobin, J. Inst. Met., 96, (1968), 236.

5. J. D. Eshelby, Proc . Roy. Soc. Lond. , 197 (1949) 396.

6. J. S. Koehler, in Imperfections in Nearly Perfect Crystals, Wiley, New

York, 1952, p. 197.

7. J. Friedel, Philos. Mag., 44 (1953) 444.

8. A. Granato and K. Liicke , J. Appl. Phys., 2J7 (1956) 789.

9. J. S. Koehler and G. deWit, Phys. Rev., 116 (1959) 1121.

10. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl. Phys., 51. (1980) 305.

11. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics, 14 (1974)

246.

12. Y. P. Varshni, Phys. Rev. B, 10 (1970) 3952.

13. H. M. Ledbetter, Phys. Stat. Solidi (a), 66 (1981) 477.

14. H. M. Ledbetter and M. W. Austin, submitted to Physica Status Solidi.

15. H. M. Ledbetter, submitted to Journal of Applied Physics.

16. W. Koster and H. Franz, Metall. Rev., 6 (1961) 1.

145
TABLE 1. Engineering elastic constants at 295 K

Reference

Copper

Copper

CDA104

Copper

CDA155

53* 80 88 63 88 96

-•-.A. 96

Ex 129 128 124 124 125 126 125 127

e 2 128 123 121 125 123 111 113

e 3 128 121 115 125 117 122 125

G 2 3 52.4

47.9 48.4 48.1 44.1 48.1 54.5 51.0

^3 2 44.9

48.3 47.5 48.8 48.1 51.9 43.7

G i 3 42.7

48.0 46.9 43.6 48.5 48.1 41.8

^3 1 42.5

48.1 46.6 45.8 48.6 47.1 41.5

Gl 2 40.0

48.4 46.3 44.8 48.3 45.5 39.1

^2 1 47.2

48.4 46.3 45.0 48.2 45.1 46.3

u 23

0.354

0.345 0.346 0.360 0.380 0.349 0.375 0.357

u 32

0.389

0.348 0.354 0.361 0.349 0.358 0.393

u 13

0.331

0.346 0.359 0.374 0.350 0.369 0.334

y 31

0.325

0.346 0.350 0.345 0.350 0.343 0.326

u 1 2 0.385

0.348 0.349 0.345 0.349 0.344 0.389

v 2 x 0.343

0.347 0.347 0.335 0.350 0.334 0.346

E = Young modulus (GPa)

G = shear modulus (GPa)

v = Poisson ratio

1 = rolling direction

3 = direction normal to rolling plane

'f = hardness (Rockwell F) = remeasured after 90 days

146
TABLE 2. Engineering elastic constants at 4 K

Reference

Copper

Copper

CDA104

Copper

CDA155

53* 80 88 63 88 96 96

Ei 140 146 142 143 143 144 142 146

E 2 146 141 139 143 140 127 146

e 3 146 139 133 143 134 139 146

^2 3 57.1

52.4 52.6 52.4 48.1 52.4 59.4 55.6

G 3 2 49.0

52.6 51.7 53.2 52.4 56.6 47.6

Gi 3 46.6

52.3 51.0 47.5 52.9 52.5 45.6

G 3 i 46.3

52.3 50.8 49.8 53.0 51.4 45.3

Gl 2 43.6

52.7 50.4 48.8 52.6 49.7 42.6

G 2 x 51.5

52.7 50.4 49.0 52.6 49.2 50.5

^2 3 0.324

0.336 0.324 0.337 0.354 0.329 0.351 0.336

^3 2 0.325

0.325 0.331 0.338 0.328 0.336 0.367

*13

0.324

0.324 0.336 0.349 0.329 0.346 0.317

*3 1 0.324

0.324 0.329 0.324 0.329 0.323 0.311

*12

0.326

0.326 0.328 0.325 0.328 0.325 0.364

*21

0.325

0.325 0.326 0.316 0.329 0.316 0.326

E = Young modulus (GPa)

G = shear modulus (GPa)

v = Poisson ratio

1 = rolling direction

3 = direction normal to rolling plane

= hardness (Rockwell F) = remeasured after 90 days

147
* TABLE 3. Fitting parameters for Varshni equation

Reference Copper CDA104 -88 CDA155-96

G G C, G

C(0) (GPa) 213.4 51.7 208.0 49.4 207.2 215.6

s (GPa) 7.42 2.41 6.82 2.24 7.97 1.97

t (K) 161.5 135.1 152.2 129.5 179.9 115.9

See Eq . (6)

Fig. 1. Microstructure of CDA104 alloy in Rp = 88 hardness condition.

View of x 2 plane: x x is horizontal and x 3 is vertical.

148
INTENSITY

1.10

TEMPERATURE (K)

Fig. 2. Temperature behavior of two typical elastic - stiffness constants:

C 22 and G 21 . The behavior is regular. For reference copper, the

intercepts at T = 0 K are 1.084 and 1.049. Thus, deformed copper

and reference copper show essentially identical C(T) behavior.

These results represent CDA104 copper in the hardness-80 condition.

i. x3» x2 x3> xi

J L x2

,x 3

x1> x 3

o o

CM T— o

]£! L*L

65 90

CM

CM

CM 115

CO ^

Standard

140

20

Fig. 3. X-ray-diffraction patterns for CDA104 alloy in Rp = 88 hardness

condition. Top two patterns represent x 3 face, rotated around x 2 and rotated around x x

. Third and fourth patterns represent x 2 and

x x faces, both rotated around x 3

. At the bottom, we show a standard pattern for annealed isotropic copper.

149
Fig. 4. An ideal block containing a single edge dislocation showing

propagation directions and polarization directions for elastic

shear waves corresponding to C 232 3 and C 2332 . For the six

principal elastic constants, the dislocation breaks usual macro- scopic s y Dime tries, c aus ing ^1111 = ^3333 ^ ^2222 and ^1221 = ^1313 = C 1331 = ^2323 > ^1212 = C 2332 .

150


ACOUSTOELASTIC RESIDUAL- STRESS MEASUREMENTS:

ROLE OF ANISOTROPIC DISLOCATION ARRAY

H. M. Ledbetter

Fracture and Deformation Division

Institute for Materials Science and Engineering

National Bureau of Standards

Boulder, Colorado

We consider how an anisotropic dislocation array affects

sound velocities and elastic constants. Assuming isotropic dislocation arrays, existing models predict that dislocations reduce

elastic stiffness up to approximately 10 percent. Measurements

show larger reductions. For an anisotropic dislocation array,

where the Burgers vectors do not average to zero over the material

volume, elastic softening is anisotropic. This means nonsymmetry

in the effective fourth-order elastic-stiffness tensor: Cjj^l

^ ^ijlk- And, in general, Cjjkl ^ . We studied this effect

experimentally by measuring sound velocities in deformed (rolled)

copper plate. We estimated sound-velocity measurement uncertainty

as less than 0.2 percent. We found that C 232 3 exceeds C 2332 by

17 percent, C 3131 approximately equals C 3113 , and C 2121 exceeds

C 2 112 by 18 percent. Such large differences indicate an enormous fictitious internal strain or residual stress when measured

acoustoelastically .

INTRODUCTION

Among various methods used to measure internal strain or residual

stress, acoustoelasticity ranks after x-ray diffraction. For an isotropic

medium, this method depends on the sound-velocity dependence on the stress

direction: a shear wave vibrating parallel to the stress axis travels with

a velocity different from a wave vibrating perpendicular to the stress axis.

The situation is equivalent to shear waves in a stress -free medium containing elastic anisotropy. Such anisotropy occurs, for example, in monocrystals and in textured polycrystalline aggregates. (By texture, we mean a nonrandom distribution of crystallite orientations.) Thus, any internal

irregularity that causes effective elastic-constant anisotropy appears as an internal stress.

*Intended for Journal of Applied Physics .

151
To determine internal stress acoustoelastically , one usually uses the

approximate relationship: 1

a . — a . i J

_ v — V

8G 2 ki ki . 4G + n v,k

( 1 )

Here, denotes the stress component along the i axis, G denotes the usual

isotropic -symmetry shear modulus, n denotes one of the Murnaghan isotropicsymmetry third-order elastic-stiffness constants (n = 4C 456 ), v denotes

acoustic -wave velocity, subscript k denotes the shear-wave propagation

direction, and subscripts i and j denote the shear-wave polarization (vibration) direction. For copper, 2 G equals 48.2 GPa and n equals —49.1. The

exact value of n remains uncertain. Here, we take the arithmetic average of

the Voigt and Reuss bounds computed from the monocrystal Cijklmn by Chang

and Graham. 3 Thus, for copper

tfi — = 129.3 (Av/v) GPa. (2)

A stress of 350 MPa, an upper-bound yield strength for work-hardened copper,

produces a Av/v of only 2.7-10-

3

. To achieve the sometimes -quoted ± 10 MPa

within 10 percent requires a Av/v measurement error less than 0.8- 10 6

!

Texture causes a large velocity difference. The largest possible

velocity difference depends on the Zener elastic-anisotropy ratio:

A = 2C 44 /(C 11 - C 12 ). (3)

For copper, with C X1 = 169.6, C 12 = 122.4, and C 44 = 75.4 GPa, we obtain

vmax/vmin = A

2 = 1.79. (4)

In practice, lower ratios occur. For example, for <100>, <11 0> , and <1 1 1>

rod textures in copper, Gairola4 found vmax/vm^ n = 1.32, 1.07, and 1.22,

respectively. If we take the lowest value, that for the <110> case, and

assume only 1 percent texture, then Eq. (1) predicts a fictitious internal

stress of 90 MPa! Thus, for acoustoelastic internal-stress measurements,

texture remains a large, unsolved problem.

This study considers another large, unsolved problem: anisotropic

dislocation arrays. We know that many engineering materials contain large

numbers of dislocations. Hull and Bacon5 estimated that well-annealed

metals contain a dislocation density of 10 6 to 10 8 cm

-2 and that heavily

deformed metals contain up to 5-10 13 cm

-2

. We expect dislocations to lower

the elastic stiffness, C, by contributing an extra strain to the usual

elastic strain:

C a

e

£ +

d

€

(5)

152
This extra strain arises because, when stressed, dislocations bow out in

their glide plane. Several authors 6

' 9 gave models for dislocation- induced

e lastic - stiffness softening. Several experimental studies 10 ' 13 confirm the

effect. For slightly deformed copper, Koehler and deWit 14 estimated that

the elastic-softening might reach 10 percent. All the models depend on non- pinned dislocations. For well-pinned dislocations, in Eq. (5) equals

zero and no elastic softening occurs. Implicitly, all the models assume an isotropic dislocation array: dislocation Burgers vectors occur in all

directions with equal probability.

This study goes a step further: it invokes an anisotropic dislocation

array. One imagines such arrays to arise usually during plastic deformation. Mughrabi 15 reviewed the topic of dislocation structures that arise

in copper during low- temperature unidirectional deformation. He described

many anisotropic dislocation arrangements: bundles, grids, and threedimensional cells. Their usually rectilinear geometry relates to the co- ordinates of the externally applied deformation forces.

To confirm the basic idea that anisotropic dislocation arrays cause

strong shear-wave-velocity birefringence, we measured sound velocities in

a copper deformed heavily by rolling. By ultrasonic methods, we measured

sound velocities corresponding to the six principal elastic stiffnesses:

C llt C 22 , C 33 , c 44 , C 55 , and C 66 . We especially focused on the differences

between ^2323 3.nd ^2332 > ^ 1

1

£ind U ^ 3 3 ^ , and ^1212 3. rid ^1221 * For a wellbehaved material, even one with strong texture, these C^j^i are equal in

pairs , that is

c ijkl = ^ijlk- (6)

As described below, this relationship fails for the studied material.

MEASUREMENTS

Material

For study, we chose a commercial copper alloy: CDA155. Including

silver, this alloy contains 99.75 percent copper. We studied three plasticdeformation conditions: the received deformed plate (hardness = 88 Rp) , a plate deformed by rolling to a 60 percent reduction (hardness = 96 Rp)

and

the recrystallized-annealed condition (hardness = 63 Rp)

From these, we ground parallelepiped specimens measuring approximately 1.5 cm for all three

dimensions. Figure 1 shows a typical microstructure. We obtained the

annealed condition by heating the received plate for 1.3 h at 550°C. We

adopted the following co-ordinate system: x x = rolling direction, x 3 = normal to rolling plane, x2 = third orthogonal direction.

153
Methods

Mass density, p, we measured by Archimedes's method using distilled

water as a standard. We estimated inaccuracy as 0.05 percent.

Sound velocity, v, we measured by a MHz -frequency pulse- echo - superposition method described previously. 16 From mass density and sound

velocities, we computed elastic constants by the general relationship

c ijij = Pvij •

Here, subscripts i and j denote wave -propagation and wave -polarization

directions, respectively. For four states of the material, we measured six

sound velocities corresponding to the longitudinal elastic constants C 11(

C 22 , and C 33 and the transverse, or shear, elastic constants C 44 , C 55 , and

C 66 . Because they contribute little to this study and because they contain

larger errors, we omitted measurements of C 12 , C 13 , and C 23 . For sound

velocities, we estimated a 0.2 percent inaccuracy in v^ and a 0.4 percent

inaccuracy in vt

. Thus, we estimated errors as 0.45 percent and C t

errors as 0.85 percent. (Subscripts Z and t denote longitudinal and

transverse waves .

)

Texture we studied by x-ray diffraction from surfaces perpendicular

to the three principal directions: x x = rolling direction, x 3 = direction

normal to rolling plane, x 2 = direction orthogonal to x x and x3 (some call

x2 the long- transverse direction). For x-ray diffraction, we used a computer-controlled commercial diffractometer and Cu-Kq radiation.

RESULTS

The study's principal results appear in Table I. The table shows

measurements for three plastic-deformation conditions. The lowest hardness

represents the recrystallized-annealed condition. For the high-hardness

condition, we show measurements at two post-deformation times: 15 days and

105 days. For brevity, we report only sound velocities and not the

The velocities correspond more closely to measurement and to any acoustoelasticity study.

Table I also shows results calculated from copper's monocrystal elastic

constants: C X1 = 168.4, C 12 = 121.4, and C 44 = 75.39 GPa. We calculated

the quasi- isotropic elastic constants by Kroner's method. 17 Figure 2 shows

a typical set of four x-ray-diffraction patterns: standard copper and patterns from the x x , x 2 , x3 faces. Study of these patterns led to the following interpretation. With respect to annealed copper, diffraction lines are broadened but little shifted. For the x x and x 2 faces, one sees essentially

identical patterns. These differ dramatically from the x3 -surface pattern.

It shows no (111) line and a strong (220) line. Unambiguously, this shows

that {110} corresponds to the rolling plane. We expect a rolling direction

of either <1 1 1> or <112>. Diffraction patterns from the x x and x 2 surfaces

suggest a mixture of both, favoring <1 1 1>

.

154
DISCUSSION

As shown by the x-ray diffraction results, the deformed material con- tains texture. Elsewhere, 18 we consider the texture effect and show that it

is small compared with dislocation effects.

As mentioned in the introduction, previous studies show large elasticconstant decreases caused by plastic deformation. For example, van den

Beukel and Brouwer 12 reported that deforming copper at 78 K decreases the

shear modulus 18 percent. It remains to determine the maximum decrease.

Since elastic-stiffness decreases are well documented and well understood, in this study we focus on symmetry breaking caused by anisotropic

dislocation arrays. Table I provides ample evidence of this phenomenon.

We see that C 1X * C 22 ^ C 33 is not explained by texture. 18 More surprisingly, we see that C 3232 C 3223 and C 1212 * C 1221 . (For another copper

material, reported elsewhere, 18 we also found that C 1313 ^ C 1331 .)

For a better understanding of the symmetry-breaking effect, consider

the usual equation of motion:

~P

3

2 u

J + 3t 2- d .a . . 1 ij = 0 . ( 8 )

Here, p denotes mass density; uj , displacement; t, time; and , the stress

tensor. Separate the displacement into elastic and dislocation parts:

u

j - u! + UJ

• (5)

and write Hooke's law:

a ij = c ijki 5 ku1 = c ijki 5 k ( ui ~ u$>- (10)

For small dislocation motions,

^ku$ ” tkimn ^mun> (11)

where amounts to an elastic polarizability. 19 Combining Eqs . (8),

(10), and (11), one obtains

~P

J- + A,. .. 0 3.3, u. = 0,

at' ijki i k I

( 12 )

where

Aijki Cijmn tmnkj>

We see that A^j^, the effective elastic-stiffness tensor, is symmetric

in i and j , the indices related to the stress tensor , but A-jj^j I s nonsymmetric in k and i

( 13 )
In Fig. 3, the two shear-wave velocities correspond to

v2 = w

2 /k2 = A3232 /

P

(14)

and

w

2 /k2

L 3 2 2 3 /P • (15)

Experiment shows that A323 2 * A3223 We understand this because in the

first case the dislocation moves through the lattice and contributes a strain corresponding to the displacement u2

. In the second case, when the

shear wave passes, the dislocation moves only with the lattice, uf = 0

.

Because A^j^i = Aji^q, we can write these two effective elastic moduli as A2332 and “2

3

2 3

# In general, we conclude that

Aijk^ * Akiij • (16)

The present study relates also to ideas of Laval 20 questioning Cauchy's

classical concept of a symmetrical stress tensor: ct-jj = crji- Laval

suggested that in dynamic elastic-constant measurements, where the passing

ultrasonic wave might cause torques and rotations, the stress tensor is

nonsymmetric . If so, then

c ijkl * cjikl- ( 17 )

Thus, a triclinic crystal, with no point-group symmetry, would exhibit

forty-five rather than twenty-one independent elastic constants. Raman and

Viswanathan2

1

studied the wave -propagation relationships and concluded that

orthorhombic symmetry would exhibit fifteen, rather than nine, independent

elastic constants. In Hooke's law, the strain tensor is mathematically the same as the stress tensor. Thus, an orthorhombic - symmetry material with a displacement described by Eq . (8) would also exhibit fifteen independent

elastic constants. Although subsequent theoretical 2

2

and experimental 2

3

studies disprove Laval's hypothesis, his symmetry ideas relate to this

study, where dislocations contribute anisotropically to displacement and

displacement gradient.

Finally, we note that acoustoelasticians have allowed formally for

permanent plastic strain. In reviewing this subject, Pao , Sachse, and

Fukuoka24 wrote the following relationship corresponding to Eq. (1) [their

Eq. (124)]:

V

ki 4G + n

+

8G 2 (18)

Here, e^— denotes the difference in principal plastic strains. Pao et.

al. call this effect acoustoelastic-plastic birefringence and point out that

the "relative change of wave speed is of the order of the permanent plastic

strain (residual strain), which can be very large." The present study

156
identifies a mechanism that can contribute to e? - e? an amount far exceeding the maximum residual strain, which corresponds to the work-hardened

yield strength.

CONCLUSIONS

From this study, we draw five conclusions:

1. Copper deformed heavily by rolling shows large elastic-constant changes.

Texture, which increases some elastic stiffnesses and decreases others,

is not the principal cause.

2. In agreement with many previous studies, both experimental and theoretical, the nonpinned dislocation density lowers the elastic stiffnesses.

3. C u , C 22 , and C 33 differ in a way not explained by texture. More

important, the macroscopic symmetry of the shear moduli C. .. .(i i= j)

breaks. We ascribe these results to anisotropic dislocation ^arrays

4. At ambient temperatures, after deformation, all the elastic stiffness

values increase with time. As in previous studies, we ascribe this

recovery to dislocation pinning by point defects.

5. For acoustoelastic studies of internal strain (residual stress), the

specter of anisotropic dislocation arrays hangs heavily. Even a small

dislocation-array anisotropy suggests large fictitious internal strains.

Clearly, such arrays exist in deformed materials. It remains to be seen whether they exist also in annealed materials.

ACKNOWLEDGMENT

Partial support for this study came from the U.S. DoE Office of Fusion

Energy. S. Kim made the sound-velocity measurements. L. Scull and R. Walsh

provided samples. Both Professor E. Kroner (Stuttgart) and Professor S. K.

Datta (Boulder) made many valuable comments.

REFERENCES

1. R. E. Green, Ultrasonic Investigation of Mechanical Properties

(Academic, New York, 1973), p. 131.

2. H. M. Ledbetter, J. Phys. D: Appl . Phys . L3, 1879 (1980).

3. R. Chang and L. J. Graham, Mater. Res. Bull. 3_, 745 (1968).

4. B. K. D. Gairola, in Continuum Models of Discrete Systems (U. Waterloo

Press, Waterloo, 1980), p. 387.

5. D. Hull and D. J. Bacon, Introduction to Dislocations (Pergamon, Oxford,

1984), p . 22

6. J. D. Eshelby, Proc . Roy. Soc . Lond. 197, 396 (1949).

157
7. J. S. Koehler, in Imperfections in Nearly Perfect Crystals (Wiley, New

York, 1952), p. 197.

8. J. Friedel, Philos. Mag. 44, 444 (1953).

9. A. Granato and K. Lucke , J. Appl. Phys . 27, 789 (1956).

10. D. 0. Thompson and D. K. Holmes, J. Appl. Phys. 27,, 191 (1956).

11. W. Lems, Physica 28., 445 (1962).

12. A. van den Beukel and C. Brouwer, Philos. Mag. L7, 453 (1968).

13. G. Fantozzi, D. Boulanger, and P. Gobin, J. Inst. Met. 96., 236 (1968).

14. J. S. Koehler and G. deWit, Phys. Rev. 116 . 1121 (1959).

15. H. Mughrabi, in Constitutive Equations in Plasticity (MIT Press,

Cambridge, 1975), p. 199.

16. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl. Phys. 51 . 305 (1980).

17. E. Kroner, Z. Phys. 151 . 504 (1958).

18. H. M. Ledbetter, submitted to Phys. Status. Solidi.

19. E. Kroner, in Theory of Crystal Defects (Academic, New York, 1966),

p. 215.

20. J. Laval, Compt. Rend. 232 . 1947 (1951).

21. C. V. Raman and K. S. Viswanathan, Proc . Indian Acad. Sci. 42, 51

(1955).

22. M. Lax, in Lattice Dynamics (Pergamon, Oxford, 1965), p. 583.

23. K. D. Schwartz and A. V. Granato, J. Acoust. Soc. Amer. .38, 824 (1965).

24. Y-H. Pao , W. Sachse, and H. Fukuoka, Physical Acoustics 17 . 61 (1984).

TABLE I. Sound velocities in units of cm/fxs.

Hardness (Rockwell F) 63 88 96 - l a 96 -2b Annealed

Pure Cu

Mass Density (g/cm3

) 8.897 8.895 8.893 8.893 8.935

v n 0.4750 0.4786 0.4787 0.4803 0.4724

v2 2 0.4753 0.4738 0.4596 0.4613

v3 3 0.4751 0.4664 0.4751 0.4769

v32 0.2324 0.2415 0.2216 0.2246 0.2318

v23 0.2324 0.2474 0.2395 0.2426

V3 1 0.2336 0.2302 0.2161 0.2186

Via 0.2335 0.2326 0.2169 0.2192

Vl2 0.2329 0.2263 0.2096 0.2121

V2 1 0.2328 0.2253 0.2281 0.2303

a Measured 15 days after deformation,

b Measured 105 days after deformation.

158
Fig. 1. Microstructure of heavily deformed

RF

96 )- View is x t plane, or along

direction is x3 , the through-plate

copper plate (hardness equals

rolling direction. Vertical

direction.

159
INTENSITY

*3,x2

JV.

J lo O o o CM T— o CM 1“ L«2_ 40 65 90

CM

CM

CM 115

2®

x2,x 3

X^Xg

o^

Standard

140

Fig. 2. X-ray diffraction patterns for deformed copper (hardness equals

Rp88) . Top two patterns represent x3 face, rotated around x 2 and

rotated around x x

. Third and fourth patterns represent x2 and x x faces, both rotated around x 3

. At the bottom, we show a standard

pattern for annealed isotropic copper.

160
Fig. 3 . Schematic to illustrate a dislocation's role in breaking elastic-constant symmetry. Under stress, the dislocation can bow in the glide plane, but not out of

the glide plane. The anisotropic reversible plastic

displacement and strain causes C 11X1 = C3333 > C 2222 . For the shear elastic- stiffness constants, it causes

^2323 = C 1313 = C 1331 = C 1221 > C 2332 = C 1212 .

161



To be submitted to A1AA Journal

Temperature Variation of the Elastic Constants of Aluminum Alloy

2090-T81

J. Glazer and J.W. Morris, Jr. Center for Advanced Materials,

Lawrence Berkeley Laboratory, and

Department of Materials Science and Mineral Engineering,

University of California, Berkeley,

Berkeley, California

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

Fracture and Deformation Division,

Institute for Materials Science and Engineering,

National Bureau of Standards,

Boulder, Colorado

Aluminum is often the material of choice for weight-critical structures used at cryogenic temperatures. Current aerospace applications include the external tank of the

space shuttle, currently manufactured from aluminum alloy 2219. Future applications

might include tanks for proposed hypersonic vehicles. Aluminum-lithium alloys have been

proposed for these applications because they provide mechanical properties comparable or superior to those of existing aerospace aluminum alloys at 7-10% lower density and higher

stiffness. Since stiffness is an important design criterion for structures like tanks, the

elastic constants for these materials at low temperature are important design properties.

This note focuses on a particular aluminum-lithium alloy, 2090-T81, that may see cryogenic service.

Aluminum-lithium alloys of commercial compositions have elastic moduli at room

temperature approximately 7-12% higher than those of conventional aluminum alloys. The

increase in elastic modulus is related primarily to the amount of lithium in the alloy, so it can vary significantly even within the specified composition ranges of commercial alloys

[ 1 , 2].

The temperature variation of the elastic constants of various commercial aluminum

alloys has been studied previously [3,4]. All aluminum alloys behave similarly, displaying an increase in elastic modulus of about 12% between room temperature and 4 K. This note

confirms this trend for the aluminum-lithium alloy 2090-T81. Titanium alloys, which are also used for cryogenic tanks, show a smaller increase in stiffness at low temperatures [4],

The alloy studied in this investigation, 2090-T81, has a nominal chemical

composition of Al-2.7Cu-2.2Li-0.12Zr in weight percent. The chemical composition

limits and the actual chemical composition, as determined by atomic absorption

spectroscopy, are given in Table 1. This alloy shows improved strength, elongation, and

fracture toughness at low temperatures. The mechanical properties at 298 K, 77 K, and

4 K are given in reference 5.

The elastic-constant measurements were performed using ultrasonic (10 MHz)

pulse techniques. The experimental procedure is described in detail elsewhere [4,6].

Except for high-strain cases, dynamic elastic constants equal static elastic constants within

the usual uncertainty of the latter. For the dynamic values, we estimate the uncertainty as 0.1%. Static and dynamic values should show essentially identical temperature behavior.

163
Figure 1 shows the Young's modulus and Poisson's ratio of 2090-T81 alloy as a function of temperature between 295 K and 4 K. The values at selected temperatures are

listed in Table 2. The room-temperature value of the Young's modulus lies close to the

reported average values from static tensile tests [2]. The temperature variation is similar to that observed for other aluminum alloys [4].

Acknowledgment: The authors thank the Aluminum Company of America for supplying

the material used in these measurements. This study was supported at the Lawrence

Berkeley Laboratory by the Director, Office of Energy Research, Office of Basic Energy

Science, Material Sciences Division of the U.S. Department of Energy under Contract No.

DE-AC03-76SF00098.

References

1. Muller, W., Bubeck, E., and Gerold, V., "Elastic Constants of Al-Li Solid Solutions

and 6’ Precipitates," Aluminium-Lithium Alloys III, eds. C. Baker, P.J. Gregson, S.J. Harris, and C.J. Peel, The Institute of Metals, London, 1986, pp. 435-441.

2. Bretz, P.E. and Sawtell, R.R., "'Alithalite' Alloys: Progress, Products and

Properties," Aluminium-Lithium Alloys III, eds. C. Baker, P.J. Gregson, S.J. Harris,

and C.J. Peel, The Institute of Metals, London, 1986, pp. 47-56.

3. Read, D.T. and Ledbetter, H.M., "Temperature Dependences of the Elastic Constants

of Precipitation-Hardened Aluminum Alloys 2014 and 2219," Journal of Engineering

Materials and Technology, Vol. 99, 1977, pp. 181-184.

4. Naimon, E.R., Ledbetter, H.M., and Weston, W.F., "Low-Temperature Elastic

Properties of Four Wrought and Annealed Aluminum Alloys," Journal of Materials

Science, Vol. 10, 1975, pp. 1309-1316.

5. Glazer, J., Verzasconi, S.L., Dalder, E.N.C., Yu, W., Emigh, R.A., Ritchie, R.O.,

and Morris, Jr., J.W., "Cryogenic Mechanical Properties of Al-Cu-Li-Zr Alloy 2090,"

Advances in Cryogenic Engineering, Vol. 32, 1986, pp. 397-404.

6. Ledbetter, H.M., Frederick, N.V., and Austin, M.W., "Elastic-constant Variability in

Stainless-Steel 304," Journal of Applied Physics, Vol. 51, 1980, pp. 305-309.

164
Table 1 Chemical composition limits for 2090 alloy and actual composition

of the material used in this study (in weight percent)

element A1 Cu Li Zr Fe Si Mg Mn Ti

composition

limits bal 2. 4-3.0 1.9-2.

0.08-0.15 0.10 0.12 0.25 0.05 0.15

actual composition bal 2.86 2.05 0.12 0.02 < 0.01 < 0.01 < 0.005 0.02

Table 2 Elastic constants of 2090-T81 alloy

Elastic

Temperature

constant

295 K 77 K 20 K 4 K

E (GPa) 78.3 86.9 87.6 87.6

(10 6

psi) 11.36 12.60 12.70 12.70

V 0.320 0.308 0.307 0.307

165
1.2

H

O

in

CD

CSJ, O

0.9

100 200 300

T (K)

Fig. 1 Young's modulus (E) and Poisson's ratio (v) of 2090-T81 alloy as functions of

temperature. The plotted values are ratios between the actual values and the roomtemperature values given in Table 2.

166
167

WELDING PROGRAM

LEADER: T. A. Siewert

CONTRIBUTORS: C. N. McCowan, D. P. Vigliotti

OBJECTIVES

• Investigation of the metallurgical factors that affect the mechanical

properties of stainless-steel weldments at cryogenic temperatures

• Contribution to the development of improved filler metals for welding

stainless steels for use at 4 K

• Evaluation of the mechanical properties of weldments at 4 K

• Development of methods for detecting defects and evaluation of their

significance in stainless steel weldments

RESEARCH HIGHLIGHTS

• High toughness, 262 MPa-m^, at 4 K was achieved in a fully austenitic

shielded-metal-arc weld that contained exceptionally high nickel content

(Fe—6Cr—43Ni) . However, its 559-MPa yield strength was relatively low.

This indicates that the improvement in the strength—toughness combination, which had been observed previously when the nickel content was

increased to 20 wt.%, may not extend beyond this value.

• Electron beam and laser welds in 25-mm-thick base metal produced 4-K

fracture toughness values near 150 MPa-mh, slightly less than the 170

to 200 MPa-mh values measured in previous tests of the AISI type 316LN

base metal. Transverse tensile specimens for both the electron beam and

laser welds fractured in the base metal, indicating that the weld did not

reduce the joint efficiency. Specimens had ductile dimple densities near

1.5 x 10 4 dimples/mm2

, a value higher than that for the base metal but

lower than that for conventional welds.

• The effect of inclusion spacing on the strength—toughness combination was described by an equation that is applicable to both weld and base metal.

The equation was developed by a regression technique that chose the best

combination of twenty variables in a large data base. Small spacing between inclusions results in a poor strength—toughness combination.

• Molybdenum and nickel interact in determining the strength and CVN absorbed energy of welds. Welds with 9- and 14-wt.% nickel and various

molybdenum contents were studied. Overall, molybdenum had a strengthening coefficient of 33 MPa per wt.%, but its effect decreased to 25 MPa

per wt.% in combination with 14-wt.% nickel. Molybdenum at the 4-wt.%

level substantially degraded the toughness when the nickel content was

9 wt.%, but had no effect in combination with 14-wt.% nickel.

169
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-13Cr-20Ni-9Mn-2Nb SMAW I, II I, II, I (159,169)

VIII IT (299) VIII (29)

Fe-15Cr-12Ni-10Mn SMAW IV IV, VIII rv (453) VIII (29)

Fe-16Cr-15Ni-4Mo SMAW IV IV rv (275)

Fe-16Cr-16Ni-2Mo FCAW II III III, VI hi (155)

VI (199)

Fe-16Cr-16Ni-4Mo FCAW IV IV, VIII IV IV (275) VIII (29)

Fe-16Cr-35Ni-4Mn SMAW I I, VIII II i (159)

[330 modified] ii (299) VIII (29)

Fe-1 7Cr-9Ni-Mn-N SMAW IX IX IX (217)

Fe-17Cr-16Ni-4Mo SMAW IV IV IV (275)

Fe-18Cr-16Ni-6 . 5Mn-2Mo SMAW IX IX IX (257)

Fe-1 8Cr-16Ni-9Mn-N GMAW III, IV IV III, IV III, VI III (155)

IV (275,453)

- V (199)

Fe-18Cr-18Ni-2Mo-5Mn-N SMAW IV IV IV (453)

Fe-18Cr-20Ni-5Mn-N GMAW IX IX IX (247)

Fe-18Cr-20Ni-6Mn-0 . 3Nb GMAW IV IV IV IV (275)

Fe-18Cr-20Ni-6Mn-0.2Ti GMAW IV IV IV IV (275)

Fe-19Cr-12Ni-2Mo SMAW II II, VIII II II (315)

[316] VIII (29)

Fe-19Cr-12Ni-2Mo E'SW IV IV IV, VI IV (275,415)

[316L] VI (199)

GMAW III, IV III, IV, IV III (155,195)

VIII IV (275) VIII (29)

GTAW III, IV III, IV, IV III (155,195)

VIII IV (275) VIII (29)

SAW III, IV, III, IV, III, IV III, VI III (155,195)

V V IV (275,415)

V (233)

VI (199)

MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES: NBSIR 78-884 (I); NBSIR 79-1609 (II); NBSIR 80-1627 (III); NBSIR 81-1645 (IV); NBSIR 82-1667 (V); NBSIR 83-1690 (VI)- NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX).

170
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF WELDS, continued

Weld Alloy Welding Tensile Charpy Fracture Fatigue Crack Report Volume

[designation] Process Impact Toughness Growth Rate (page)

Fe-19Cr-12Ni-2Mo SMAW II, III, II, VIII II, III, III, IV, VI II (299,315)

[ 3 1 6 L ] . continued IV IV III (155,167)

VI (199) VIII (29)

Fe-19Cr-12Ni-2Mo-N SMAW II II, VIII II II (299)

[ 31 6LN] VIII (29)

Fe-19Cr-16Ni-7Mn GMAW IV IV IV (453)

Fe-20Cr-10Ni SMAW I I, VIII I (169)

[300] VIII (29)

Fe-20Cr-10Ni FCAW III III III, VI III (155)

[308L] VI (199)

GMAW III III III, VI III (155)

VI (199)

SMAW II, III II, VIII II, III III, VI II (299)

III (155)

VI (199) VIII (29)

Fe-20Cr-10Ni-N FCAW IV IV IV, VI IV (275,415)

( 308LN] V (199)

Fe-20Cr-10Ni-9Mn-N SMAW I I I (159,169)

Fe-20Cr-16Ni-7Mo-2W SAW IV IV IV (275 )

Fe-20Cr-34Ni-2Mo-Nb SMAW VII VII VII (253)

Fe-21Cr-6Ni-9Mn-N SMAW I, II I, II I (159,169)

II (299)

AUSTENITIC STEELS

Fe-1 5Mn-8Ni-lMo-0 . 7C SAW VIII VIII VIII (15)

Fe-5Cr-25Mn-lNi V V.VII V,VI V (233)

VI (199) VII (245)

Fe-6 . 5Cr-25Mn-3Ni-lMo SMAW V V V,VI V (233,245)

VI (199)

ALUMINUM ALLOYS

Al-5Mg GMAW III IV IV III (155,217)

[5183] IV (323)

Al-5Mg-0.12Ti GMAW III IV IV III (155, 217)

[5556] IV (323)

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

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

NBSIR 84-3000 (VII); NBSIR 85-3025 (VIII); NBSIR 86-3050 (IX).

171



THE ROLE OF INCLUSIONS IN THE FRACTURE

OF AUSTENITIC STAINLESS STEEL WELDS AT 4 it

T. A. Siewert and C. N. McCowan

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

Inclusion densities were measured for three types of

austenitic stainless steel welds and compared to the 4-K yield

strengths, 76 -K Charpy V-notch absorbed energies, and the ductile

dimple densities on the respective fracture surfaces. The welds

included shielded metal arc (SMA) welds and gas metal arc (GMA)

welds. The inclusion density was consistently a factor of 8 to 10

less than the fracture surface dimple density. Inclusion and dimple densities ranged from 3.9 x 10 4 inclusions -mm" 2 and 3.2 x 10 5 dimples -mm" 2 for one SMA specimen to 1.1 x 10 4 inclusions • mm" 2 and 1.2 x 10 5 dimples -mm" 2 for the fully austenitic GMA specimen.

Both ductile dimple density and dimple morphology varied with

specimen type. The inclusion data for the welds agreed with a linear relationship between fracture toughness and inclusion

spacing that had been developed for base metals.

INTRODUCTION

Type 316LN stainless steel is emerging as the preferred structural

material for superconducting magnet casings designed to operate at 4 K

because it has a higher strength than the older 316L grade and comparable

toughness at cryogenic temperatures. 1 ’ 3 Unfortunately, welding electrodes

matching this composition and strength have substantially poorer toughness

at 4 K than the stainless steel plate. This is illustrated in figure 1 for

type 304LN stainless steel (similar in strength and toughness to type 316LN)

and various welds produced with type 308LN and type 316LN electrodes. 4 Matching the electrode strength to the base metal strength results in much

lower toughness. This forces the designer to move the welds into regions

of the structure where the lower toughness can be tolerated. Where high

toughness welds are required, the corresponding weld strength is substantially reduced, and the structural design becomes less efficient. This

study sought the reason for the lower weld toughness and also looked at

parameters that can be used to predict the 4-K fracture toughness.

^Submitted to Advances in Cryogenic Engineering—Materials . vol. 34.

173
EXPERIMENTAL PROCEDURES

Two shielded metal arc (SMA) welds and a gas metal arc (GMA) weld were

selected for detailed microstructural evaluation. These welds were specifically selected on the basis of previous mechanical property evaluations, so that their range of toughness and strength would reveal any microstructural

basis for variations in toughness. 5 ’ 6 Table 1 lists the mechanical property

data and composition of these welds, Welds SMA- 1 and SMA- 2 1 differ primarily

in nitrogen and manganese content. Data from previous studies indicate a reduction in Charpy V-notch (CVN) absorbed energy as nitrogen additions

increase the strength. 6 Because GMA has twice the nickel content of SMA- 21,

the importance of the nickel content in increasing the CVN absorbed energy

and its effect on the fracture surface appearance could be examined. The

delta- ferrite content is also known to affect toughness, but the compositions

in this study vary over a small range (0—4 FN) , and Szumachowski and Reid

have shown that the effect of delta- ferrite content is less than that of the

nitrogen content. 7

Specimens for microstructural evaluation were selected from the 25 -mm- thick butt welds used for mechanical property tests. Fracture surfaces that

were analyzed were those of specimens fractured at 4 K, CVN specimens fractured at 76 K, and compact tension (CT) specimens fractured at 4 K. The

surfaces chosen to be polished were parallel to the direction of the welding

and normal to the plate surface, at least 1 cm in each dimension, and they

included several weld passes. They were polished until scratches were no longer visible at 3000X, but the polishing time was minimized to reduce

relief effects and selective removal of inclusions. This finish enabled

measurement of the inclusion size and density with a scanning electron

microscope (SEM) having a quantitative image analyzer.

Examination of both the fractured and polished surfaces at 3000X

enabled easy identification of the micrometer - scale inclusions characteristic of these welds. This magnification was quite useful measuring the

surface dimple dimensions and densities. At least 500 inclusions and

dimples were counted to provide statistically significant data.

The all-weld-metal specimens for tensile testing were 6 mm in diameter and oriented parallel to the direction of welding; their fracture sur- faces were, therefore, normal to the direction of welding. Tensile tests

were performed in accordance with ASTM standard A370 at a strain rate of

2 x 10' 4 s” 1

. Details of the fixture and the procedure for tensile testing

at 4 K have been described elsewhere. 8

Specimens for CVN tests were standard 10-mm square, centered in the

weld and notched so that the crack propagated in the weld direction; they

were prepared according to ASTM E23. Because physical property constraints

(adiabatic heating and low heat capacity) preclude obtaining meaningful data

at 4 K, the CVN testing was performed at 76 K.

The 25 -mm- thick CT specimens were notched so that the fatigue crack

propagated down the center of the weld in the direction of the welding; they

were prepared according to ASTM E813. The KIc (J) value was calculated from

174
the fracture energy, J, using the single -specimen unloading- compliance

technique. Details of the fixture and testing procedure have been described

elsewhere. 9 The CT data (average of two specimens) were available for only

the GMA weld.

RESULTS AND DISCUSSION

Role of Inclusions

Neither inclusion density nor the average inclusion size was affected

by the differences in nitrogen content of the SMA-21 and SMA-1 welds (see

table 2). The SMA-21 weld had greater yield strength owing to its higher

nitrogen content, but the inverse relationship between strength and toughness observed in a previous study6 is not evident here. The similar inclusion densities and sizes of SMA-21 and SMA-1 may be the reason for the

similarity in their CVN absorbed energies. This suggests that the factors

controlling strength and toughness might be considered independently, at

least for these two welds.

The GMA weld metal has more than twice the CVN absorbed energy of

the SMA welds, yet its strength is similar to SMA-21. Therefore, the

difference in toughness is due to another factor, possibly the nickel

content, inclusion content, inclusion size, or a combination of these

factors. Unfortunately, these factors are confounded in these data.

A study by Reed and Simon separated the two factors of nickel content

and inclusion size. 3 Although their study was of base metals, the stainless

steel compositions were similar to these weld-metal compositions. Through

a statistical regression analysis of NBS test data, they determined that

the fracture toughness at 4 K could be predicted by a linear combination of

strength, nickel content, and inclusion spacing (calculated from inverting

and taking the square root of the inclusion density)

KIc (MPa-mH

) = 130 - 0 . 34a

y + 20[Ni] + 2.25A, (1)

where ay

is the yield strength (MPa), [Ni] is the nickel content (wt.%) and

A is the inclusion spacing. The standard deviation for the fit of this

equation to 19 data points was 31 MPa-m^. This equation was developed from

data with a range of 550 to 1100 MPa, a [Ni] range of 8 to 12 wt.%, and a A range of 40 to 110 fim.

The 4-K fracture toughness of the GMA weld (table 1) is predicted well

by equation 1, even though it was developed from data on alloy 316-type base

metals. Since fracture toughness data at 4 K are not available for the SMA

welds in table 1, they cannot be used to evaluate the fit of the equation.

However, 4-K fracture toughness has been evaluated in previous studies of

austenitic stainless steel welds. 10-12 Therefore, inclusion densities were

measured on three of these welds. These data, listed in table 3, also fit

the equation fairly well, although they are outside the range of composition

(Ni, Mn) of alloy 316 for which the equation was developed. This is sur- prising because welds and base metal have different cooling rates and

175
solidification structures and an order-of -magnitude difference in inclusion

density. This emphasizes the importance of inclusion density in controlling

fracture toughness. When the weld metal data are added to the base metal

data, the predictive equation for fracture toughness can be revised to cover

this expanded inclusion density range.

When revising this predictive equation, it is also appropriate to re- assess the form. Because the fundamental relationships among toughness,

strength, and composition are only beginning to be developed, it is difficult to choose the optimum form for a predictive equation. For this reason,

a stepwise multiple linear regression program was used to choose the statistically most significant form from a selection of over thirty combinations

of strength, nickel content, and inclusion spacing. The combinations included linear, product, and quotient forms with various exponents. The

revised predictive equation selected by the regression program is

KIc = 30 + 6000 ^^(1 + 0.02A), (2)

°y

where the units are the same as in equation 1. This equation, predicting

the toughness for both weld and base metal compositions, was developed from

the data used for equation 1 and four weld data with a Oy range of 550 to

1100 MPa, a [Ni] range of 8 to 20 wt.%, and a A range of 5 to 110 nm. It

has a standard deviation of 27 MPa-m^, an R

2 of 0.87, and an F ratio of 159.

This standard deviation is comparable to that found for equation 1, which

was developed for the restricted data set of alloy 316- type base metals.

Role of Dimple Density

All fracture surfaces were composed of ductile dimples. The inclusions

that nucleated the dimples were still at the bottom of some of the dimples.

Microprobe analysis indicated they were all of the manganese silicate type.

The fracture -surface dimple measurements are given in table 4. For each

weld, we see less than 25 percent variation when the dimple densities were

measured on tensile, CVN, and CT (for one weld) fracture surfaces. This

variation is small considering that testing was done at two temperatures,

at strain rates ranging from 2 x 10* 4 s* 1 for the tensile and CT specimens

to approximately 1 x 10 3 s" 1 for the CVN specimens, and that the specimen

types had different loading configurations. However, the CVN data are con- sistently higher than those for the other fracture surfaces, which may indicate a small effect. More detailed studies are needed to confirm this

trend. The good agreement found here indicates that the dimple densities

of the various welds can be characterized with any of these specimen types.

The higher CVN absorbed energy of the GMA weld could be explained by

the lower density of inclusions measured on the polished surface section.

Assuming all ductile dimples observed on fracture surfaces were nucleated

at inclusions, one can correlate the higher CVN absorbed energy to the lower

dimple density.

In a study of C—Mn welds, Passoja and Hill found very good agreement

between the inclusion density measured on an electropolished planar surface

176
and the number of dimples on a fracture surface. 13 In fact, they found that

the 1.18 factor in an equation by Kochs,

D s - 1.18 Ns

-is

, (3)

was sufficient to correct the results when the inclusion density was calculated in terms of a three-dimensional nearest-neighbor spacing. 14 Here D s is the nearest-neighbor spacing, and Ns is the average planar surface inclusion density.

Comparison of tables 2 and 4 reveals about an order-of -magnitude

difference between the inclusion density on planar surfaces and the fracture surface dimple density for this study. This difference in density is

clearly visible in figures 2 and 3. The largest difference predicted by

the fracture models for these two surfaces (corrected for the difference

between plan views of a 2- and 3 -dimensional surface) is about a factor

of 2. 16 The unexplained difference (about a factor of 5) could be due to

undercounting the inclusions on the polished surface, overcounting the

dimples on the fracture surface, or a surface roughness greater than that

predicted by general fracture models.

The first possible cause, undercounting inclusions on the polished

surface, is unlikely since the surfaces were polished with care and the

final polishing time was minimized. Pits indicating that inclusions had

been removed from the surface were not evident. The electropolishing technique employed by Passoja and Hill, may, in fact, have led to an overestimate of inclusion content in their study owing to selective removal of the

matrix material. The second possible cause, overcounting of the dimples,

is also unlikely. Whether the small indentations in figure 3 are facets of

one large inclusion or several smaller inclusions is somewhat uncertain, but

the possibility of miscounting by a factor of 5 is very small. This leaves

the roughness of the fracture surface as the most likely reason for the differences between inclusion counts and dimple density. Grain orientations

that favor shear on other than a 45 -degree angle, nonuniform inclusion distribution, and other micros true tural heterogeneity could cause increased

surface roughness.

Comparison of the inclusion sizes of the polished and fracture sur- faces provides further insight into the fracture process. Table 2 includes

average inclusion diameters measured on the polished surfaces. Figures 4

through 6 are plots of inclusion diameters measured on the fracture surfaces

of these three welds versus the average dimple diameters. In some cases,

the dimples had major and minor axes, and the data are represented with a bar. In other cases, the dimples were circular and are represented by an X within a circle. Notice that there were relatively few inclusion—ductile

dimple pairs that could be measured. Many dimples had lost their inclusions

during the fracture process. Also, notice that the average diameter of

the fracture -surface inclusions is much greater than that of the polishedsurface inclusions. This may indicate that the small inclusions were lost

more frequently than the large ones or that large inclusions are more likely

to initiate fracture and, therefore, they are more commonly observed on the

177
fracture surface. We found that another possibility, the correction to the

true inclusion diameter from random sections of a polished surface, increased the average inclusion diameter only slightly.

The measurements shown in figures 4 through 6 also indicate that larger

inclusions form larger dimples. Since the higher toughness GMA weld had

larger dimples, perhaps further improvements in toughness can be achieved by

reducing the inclusion density through better refining or agglomeration of

the manganese silicate into larger, widely spaced inclusions.

CONCLUSIONS

1. The fracture -surface dimple density and inclusion density were inversely related to the Charpy V-notch absorbed energy, indicating that

inclusion spacing is a primary factor in determining the toughness.

2. The fracture -surface dimple density was found to be almost independent

of the strain rate during fracture and specimen type. It varied less

than 25 percent when measured on a tensile specimen fractured at 4 K, a Charpy V-notch specimen fractured at 76 K, and a compact tension

specimen fractured at 4 K.

3. The fracture -surface dimple density was a factor of 5 greater than

that predicted by fracture models based on inclusion density. This

indicates either shear on an angle other than 45 degrees, a nonuniform

inclusion distribution, or microstructural inhomogeneity.

4. The weld metal fracture toughness can be predicted from an equation

developed for the base metal. The applicability of one equation for

both weld and base metal indicates a role of strength, Ni content, and

inclusion density in determining the toughness of the austenitic stainless steels.

ACKNOWLEDGMENT

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

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Processes, ASTM STP 600, American Society for Testing and Materials,

Philadelphia, 1976, pp . 30—46.

14. U. F. Kochs: Philos. Mag., 1966, vol. 13, pp. 541—544.

15. E. E. Underwood: J. Met., 1986, vol. 38, pp. 30—32.

Table 1 Selected Data for the Welds*

SMA-1 SMA-21 GMA

Yield Strength (MPa at 4 K) 460 880 1020

CVN Absorbed Energy (J at 76 K) 45 42 102

KIc (J) (MPa-mh

) — — 203

Ferrite Number (FN) 4 0.5 0

Composition (wt.%)

C 0.03 0.03 0.03

Mn 1.6 3.5 5.4

Si 0.29 0.37 0.26

Cr 17.6 17.8 18.1

Ni 9.2 9.1 20.4

N 0.05 0.15 0.16

*This study.

179
Table 2. Inclusion Measurements on Polished Surfaces

Measurement

Weld

SMA1

SMA- 21 GMA

Inclusion Density, n (mm -2

) 3.9 x 10 4 4.0 x 10 4 1.1 x 10 4

Average Diameter (/.zm

)

0.36 0.37 0.45

Volume Fraction 4.8 x 10' 3 4.9 x 10' 3 2.3 x 10' 3

Mean Spacing, A (/xm) 5.1 5.0 9.5

Table 3. SMA Fracture Toughness Data at 4 K Developed in Previous Studies

Weld

Identification

Klc

(MPa-m^)

Oy

(MPa)

Ni

(wt.%)

n

(mm' 2

)

A*

( m®) Reference

497-5 FCA 159 608 15.2 4 x 10 4 5 10,11

508-2 GMA 163 779 11.8 4.2 x 10 4 4.9 10,11

1616-2 FCA 235 607 16 3.2 x 10 4 5.6 12

A (the inclusion spacing) is defined as the inverse root of n (the

inclusion density)

.

Table 4. Dimple Measurement on the Fracture Surfaces

Specimen

Weld

SMA- 1 SMA- 21 GMA

4-K Tensile Specimen

Dimple Density (mm" 2

) 3.2 x 10 5 3.5 x 10 s 1.2 x 10 s 76-K CVN Specimen

Dimple Density (mm' 2

) 4.1 x 10 s 4.6 x 10 s 1.6 x 10 5 4-K CT Specimen

Dimple Density (mm' 2

) 1.2 x 10 s

180
Figure 1

Comparison of the strength—toughness relationships for type 304LN

austenitic stainless steel and type 308LN and 316LN welds.

181
Figure 3. Fracture surface of GMA tensile specimen.

2 4 6 8

Dimple Diameter, jum

Figure 4. Dimple diameter versus inclusion diameter for weld SMA-1.

182
2.0

1.5 ~

Inclusion

Diameter, 1.0 -

jum

0.5

i

1

m2) 1

' _h—

1

2 4 6 8

Dimple Diameter, jum

Figure 5. Dimple diameter versus inclusion diameter for weld SMA-21

Dimple Diameter, ^m

Figure 6. Dimple diameter versus inclusion diameter for GMA weld.

183



CRYOGENIC STRENGTH, TOUGHNESS, AND FERRITE CONTENT

OF STAINLESS STEEL WELDS*

Christopher N. McCowan

To study the effect of alloy content on the strength, toughness, and ferrite content of shielded metal arc (SMA) stainless

steel welds, two test matrices were evaluated.

Evaluation of a SMA weld test matrix in which manganese (1.5

to 10 wt.%) and nitrogen (0.04 to 0.26 wt.%) were varied independently has clarified the effect of these elements on the cryogenic

mechanical properties. Several molybdenum and boron additions

were also made. 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

76 -K Charpy V-notch (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 and

molybdenum slightly increased the yield strength and slightly decreased the lateral expansion. The solid- solution strengthening

of molybdenum was greater than that of manganese and the 3.84 wt.%

molybdenum alloy had substantially lower toughness than other

similar compositions. The 4-K tensile strength was relatively

unaffected by alloy additions; values varied between 1300 and

1500 MPa.

The second SMA test matrix was evaluated using stepwise

linear regression techniques to determine more accurate equations

for predicting the ferrite number (FN) in welds having high alloy

contents. This test matrix contained the weld compositions and

associated FN values of the manganese—nitrogen test matrix plus

those of approximately 850 other SMA welds.

*M.S. Thesis, No. T3449 , Publication of Colorado School of Mines,

Golden, Colorado, 1987.

185
The simplest predictive equation (FN1) developed had chromium

and nickel equivalents similar in complexity to the DeLong equivalents. The coefficients for chromium, nitrogen, and silicon, however, were lowered with respect to the DeLong coefficients while

those for carbon and nickel remained approximately the same. When

using the FN1 equation, approximately 70 percent of the welds ana- lyzed had predicted FN values within 2FN of the measured values.

When the DeLong- type equation was used, only around 50 percent of

the predicted FN values fell within this error range.

Several alternative predictive equations were also developed.

A modification of the "pseudo polar" Schaeffler equation predicted

FN with accuracy close to that of the FN1 equation. Another, more

complex, predictive equation developed included interaction terms,

such as molybdenum—chromium and manganese—nitrogen. The results

indicated that more accurate prediction of FN requires: (1) better

understanding of the interactions between compositional elements,

(2) development of nonlinear predictive equations forms, and

(3) development of separate FN predictive equations for welds with

primarily ferritic solidification modes and welds with primarily

austenitic solidification modes.

I

.

INTRODUCTION

A. Overview

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

compositions, the welds still do not match the toughness of the base material. Recently, controlled additions of nitrogen (along with additions of

manganese to reduce the tendency to form porosity) have increased the plate

strength at cryogenic temperatures with little loss in toughness. 6

Typical mechanical properties for nitrogen- strengthened base materials

are reflected by the properties reported for an as-received hot-rolled 316LN

stainless steel plate. 7 A 4-K yield strength and fracture toughness [ Kj c

( J )

]

of 1020 MPa and 224 MPa./m, respectively, were reported, with an associated

76-K CVN absorbed energy of 89 J and a lateral expansion of 1.03 mm. This

increased base materials strength, however, cannot be totally utilized until

welds of comparable toughness at these elevated strengths are developed.

To guide electrode development projects, the 1984 NBS—DOE Workshop on Materials at Low Temperatures established a 4-K fracture toughness goal of

186
160 MPa7m [KIc (J)] for weld metals having a yield strength of 1000 MPa. 8 As an economical means of screening promising alloy compositions, the workshop suggested 76-K Charpy V-notch testing.

In addition to mechanical property goals to guide electrode development, more accurate methods for predicting ferrite content in these highly

alloyed welds are needed. Unlike the fully austenitic 300-series base

materials, the welds normally have microstructures containing 3 to 7 percent

ferrite. The presence of the more brittle ferritic phase reduces toughness

at low temperatures. Small amounts of ferrite, however, are generally

desirable to reduce microfissuring tendencies in the weld that would more

severely degrade the toughness. To achieve an accurate, optimum ferrite

content of experimental alloys, new ferrite predictive equations must be

developed.

B . Objective

The test matrices evaluated in this study were designed to help clarify

the effects of nitrogen, manganese, and molybdenum on the strength and

toughness of shielded metal arc (SMA) weld metals at low temperature, and

the ferrite content of weld metals. The study was intended to contribute to

the development of welds having mechanical properties comparable to those of

the nitrogen- strengthened base materials and to develop equations that

better predict the ferrite content of these more highly alloyed weld metals.

C . Literature Review

In this review, emphasis is placed on the effect that ferrite, nitrogen, manganese, and molybdenum have on the strength and toughness of 300-

series stainless steels at low temperature and on the prediction of weld

metal ferrite content. In addition, the deformation behavior of stainless

steels during tensile testing at 4 K and the difficulties of comparing

Charpy V-notch absorbed energy to lateral expansion results for materials

of differing strengths is discussed.

1 . Strength

The strengthening potential of nitrogen in stainless steel increases

with decreasing temperature. 5 ’ 9

" 12 Nitrogen occupies the octahedral interstitial sites in the face centered cubic structure and expands the lattice

in a manner similar to that of carbon. The carbon-atom size is larger than

that of nitrogen, but when the temperature is decreased from room temperature to 4 K, interstitial nitrogen increases the yield strength by approximately sixfold compared with the twofold yield strength increase associated

with carbon.

It is clear that the traditional "hard-ball-model" 13 ’ 14 for strengthening by lattice expansion does not adequately address the differences

between carbon and nitrogen strengthening in stainless steels at low temperatures. Stronger metal-nonmetal bonding occurs in Fe-C than in Fe-N,

187
because the extra electron associated with the nitrogen atom may result in an increased elastic misfit for the nitrogen atom, which would partially

explain this difference. 15 ’ 16 The solute atom—dislocation interaction would

increase and result in an increased yield strength.

Unlike high carbon contents, high nitrogen contents do not result in

sensitization problems. Nitrogen content can be increased until porosity

occurs at the solubility limit. For a liquid Fe—18Cr—UNi alloy, solubility

is not reached until approximately 0.2 wt.% nitrogen. 17 Nitrogen solubility

increases as the chromium content in the alloy increases. The absorption of

nitrogen obeys Severts law in alloys of up to 57 wt.% chromium. 18 ’ 19 In

commercial alloys, the manganese additions also increase nitrogen solubility

so that nitrogen contents greater the 0.2 wt.% are attainable. Clearly, the

high solubility limit, lack of sensitization, and its greater strengthening

power make nitrogen an attractive interstitial alloying element.

Further increasing its attractiveness is the ability to add nitrogen to

weld metal deposits easily and economically. Nitrogen can be introduced to

the weld through shielding- gas mixtures with the gas metal arc (GMA) and gas

tungsten arc (GTA) processes or through nitrogen- containing alloy additions

to the coating with the shielded metal arc (SMA) process.

The increase in yield strength at low temperatures due to nitrogen

alloying in type-308 stainless steel composition is well known. To simplify

comparisons of studies done within differing nitrogen content ranges, a linear relationship is assumed between nitrogen and yield strength and the

effect of nitrogen is expressed as an approximate function of nitrogen content. With increases in nitrogen contents of 0.05 wt.% at 76 K, increases

in yield strengths of 150 and 175 MPa were reported by Enj o et al. 11 and

Onishi et al., 12 respectively. At 111 K, Mukai et al. reported an increase

in yield strength of 120 MPa. 20 In comparison, only slight increases in weld

metal yield strengths have been found at room temperature. 11 ’ 12 ’ 20 ’ 21

The influence of nitrogen on the tensile strength is disputed. Enjo

et al. found tensile strength increases of less than 10 MPa per 0.05 wt.%

nitrogen, implying a nearly constant tensile strength at 76 K with respect

to nitrogen content. 11 Onishi et al. found increases in the tensile

strength of approximately 25 MPa per 0.05 wt.% nitrogen. 12 In all cases,

the increases in tensile strength reported were much less than those

observed for yield strength.

The effects of manganese, molybdenum, and boron on cryogenic weld metal

strength have not been published. For 300- series -base materials, however,

some data on the effects of manganese and molybdenum are available. Manganese has been reported to increase the 76 -K yield strength slightly in

Fe—12Cr—12Ni—Mo—N alloys by Mukai et al. 20 and by Sakamoto et al. 22 in a wide range of stainless steel compositions. Yamamoto et al. 23 report that

manganese had no effect on strength, but that molybdenum increased the yield

strength at 4 K by 80 MPa for each wt.%. Purtscher and Read24 report a 40 -MPa yield strength increase per wt.% molybdenum at 4 K. Mukai et al. 20

report a 50-MPa increase in yield strength per wt.% molybdenum at 77 K.

188
2 . Toughness

The cryogenic toughness of type-308L and -316L stainless steel welds is

determined by the interrelated and sometimes opposing effects of the individual alloying elements and the ferrite content. The ferrite content is

often represented by a ferrite number (FN) , which is approximately equivalent to percent ferrite within the range of 0 to 20 FN. 25

The FN was found to be a dominant factor in determining Charpy V-notch

(CVN) toughness for these 300-series stainless steels. 26 ' 29 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 (AE)

and lateral expansion (LE) values in SMA welds. 4 Contradicting this re- ported nitrogen effect, Enjo et al. 11 found that CVN absorbed energies increased with increasing nitrogen content (decreasing FN) in his GMA study of

type- 308 stainless steel welds. Whether increasing nitrogen content lowers

the CVN toughness more than decreasing FN raises it is still unclear. Separating these factors to evaluate their individual effects on toughness is

difficult. To evaluate the effect of nitrogen, for example, FN must be held

constant by concurrently raising the chromium content or that of some other

ferritizing element. Past research indicates, however, that increased chromium and carbon contents can reduce toughness. 4 ’ 20 ’ 28 The effects of molybdenum are not clear, but Miura30 reports a dramatic drop in CVN absorbed

energy in a Fe—12Cr—12Ni—0 . IN base metal alloy for molybdenum contents of

approximately 7 to 9 wt.%. This type of loss in toughness has been reported

at molybdenum contents as low as 3.5 wt.% in 25Cr—16Ni—0 . 34N—Mo welds. 31

Manganese has been reported to decrease the 77 -K CVN absorbed energy in

stainless steel base material but to slightly increase fracture toughness at

4 K. 24 Nickel is the only compositional element common to stainless steels

that is generally agreed to increase CVN toughness.

Charpy V-notch test results for type-308L stainless steel welds at

76 K show substantial scatter. To reduce the scatter, only weld deposits

having FN less than eight and low nitrogen contents are reported here. For

SMA weld deposits, absorbed energy values ranging from 16 to 34 J and lateral expansions from 0.15 to 0.46 mm have been reported. 5 ’ 11 ’ 32 When GMA

and GTA 308L stainless steel welds are included absorbed energies of 29 to

greater than 100 J have been reported. 5 ’ 11 The principal differences in

toughness values of welds produced with different processes is attributed

to the lower inclusion content associated with the GMA and GTA welds.

3 . Ferrite—Number Prediction

The strong influence of ferrite content on cryogenic toughness dictates

the use of an accurate method for predicting the FN in the new high nitrogen and manganese stainless steel electrodes. The diagrams developed by

Schaeffler 33 and DeLong et al. 34 ’ 35 are the most common methods for the

prediction of FN (figures 1 and 2).

The basis for these diagrams can most easily be seen from the Fe—Cr-Ni

isothermal diagram at 1000°C (figure 3a). In figure 3b a portion of this

189
isothermal diagram is shown in an orientation similar to the Schaeffler and

DeLong diagrams. It is apparent that the general shape of the austeniteplus-ferrite (7 + a) two-phase region and the orientations of the singlephase (7 and a) regions with respect to it are carried over to the FN

predictive diagrams. The FN predictive diagrams are, however, nonequilibrium constitutional diagrams that were empirically developed from alloys

having more complex compositions than those of the Fe—Cr—Ni ternary diagram

alloys. These differences make further comparisons between the two types

of diagrams difficult.

Although both the Schaeffler and DeLong diagrams have proven quite

successful and useful in predicting the FN, recent studies 36 ’ 37 indicate

that several of the coefficients used with these diagrams do not accurately

reflect the appropriate trends. Also, alloys are being developed with

compositions outside the ranges of applicability for these diagrams.

The Schaeffler diagram has been found to be an accurate predictor of

FN for type-308 and -347 stainless steels. DeLong et al . 34,35 refined the

Schaeffler diagram to better predict FN for type-309 and -316 stainless

steel compositions. These compositions have higher chromium and nickel

equivalents than the type-308 and -347 stainless steel welds for which the

Schaeffler diagram is best suited. Whether the respective FN diagrams or the corresponding mathematical expressions for these diagrams are used,

nickel and chromium equivalents are commonly used to represent the relative

austenitizing and ferritizing tendencies for a given alloy composition.

The Schaeffler equivalents are:

Ni eq = Ni + 30(C) + 0 . 5 (Mn) ; ( 1 )

Creq = Cr + Mo + 1.5(Si) + 0.5(Nb); (2)

where the elements are given in wt.%

DeLong added a nitrogen term with a coefficient of 30 to the Schaeffler

nickel equivalent and refined his diagram to represent only those alloy

compositions of more common commercial interest. The constant FN lines on both diagrams, however, vary in slope and spacing. Therefore, both diagrams

would best be represented by a polar- type equation form that would allow

both the slope and the spacing changes of the constant FN lines to be

modeled. The Schaeffler diagram is expressed mathematically by a "pseudo

polar" equation form : 38

Cr - 5.5

eq

FN (Schaeffler) = -39.1 + 43.5 — —- ' ( 3 ) Ni + 2 .

0

v ' eq

Presently , however , the best representation of the DeLong diagram is by an equation having a linear form : 38

FN (DeLong) = 30.64 + 3.49[Creq ] - 2 . 5 [Nieq ]

. (4)

190
This linear FN predictive equation is a reasonably accurate representation

of the DeLong diagram at low FN, but at FN values of greater than ten, the

predictive error increases.

Two important advantages of the linear representation of the DeLong

diagram and the respective chromium and nickel equivalents used within it are simplicity and clarity. Alloy elements that promote the formation of

ferrite, such as chromium, molybdenum, niobium, and silicon, are grouped

together in the chromium equivalent. Austenitizing elements, which include

nickel, carbon, nitrogen, and manganese, make up the nickel equivalent.

Each elemental variable within the respective equivalent expressions is

weighted and the difference between the equivalents determines the FN.

Chromium and molybdenum are shown to have equivalent ferrite- forming powers

on a weight percent basis. Niobium promotes ferrite formation only half as well. Nickel is approximately 28 percent less effective in preventing ferrite formation as chromium is in promoting it, and so on. The problem with

making these empirical interpretations is that all the FN predictive equations are designed around specific alloy compositions.

Thermodynamics would predict an FN predictive equation of a more

complex form: 39

FN = KNi Ni + KMn Mn + KCr Cr + % N + K* i (5) + MnNi + NiCr + K^_j ij .

This solution form implies that, in addition to the effect on ferrite

formation by individual alloying elements, interaction terms between these

elements need consideration to more accurately predict FN over a wide range

of alloy compositions. Support can be found for the need to include interaction terms in FN predictive equations. For example, the austenitizing

power of manganese has been observed to decrease when concentrations exceed

4 to 5 wt.% 40 The indication of a chromium—nickel interaction is found as far back as 1939 in the Maurer diagram. 41 Nitrogen—manganese interactions

have been suspected. 42 ’ 43

The struggle for better definition of coefficients in the FN predictive

equations is indicated by the various coefficients and constants that have

been proposed to model the effects of nitrogen and manganese (table 1 ) .44-48

At high nitrogen contents, Espy found that the nitrogen coefficient decreases. 37 Most studies involving high manganese contents have proposed

constants to represent the effect of manganese and replace the 0.5 term of

DeLong et al. 36,37 The Hull study43 was the only high manganese test matrix

to yield a manganese coefficient. However, its second- order term does re- duce 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 for proposing different terms to model the influences

of manganese or nitrogen in the nickel equivalent expression. The specifics

of any possible interactions remain undefined.

In addition to the need of defining the appropriate variables and

coefficients for better predicting FN, the ambiguity of predicting ferrite

191
formed via primary ferritic versus primary austenitic or mixed solidification modes from the same function of composition must be addressed. The

DeLong diagram, for example, was developed empirically from types 308, 309,

310, 316, 410, and 502 stainless steel alloys. These alloys solidify via a primary ferritic solidification mode. So the Delong diagram, by the nature

of its development, has empirically incorporated the kinetic factors asso- ciated with primary ferritic solidification into the positioning of the

constant FN lines. For weld deposits having FN values near zero, however,

fully austenitic or mixed mode solidification mav have been involved in

the freezing process. It might be speculated that the quantity of ferrite

formed by either of these processes could not be accurately predicted by

constant FN lines representing the primary ferritic solidification process.

For clear and accurate definition of the constant FN lines from zero to

five, it may be necessary to separate weld data by solidification mode.

Although there are several proposed mathematical expressions to predict

whether primary austenitic or ferritic solidification will occur, as yet,

these types of expressions have not been integrated into the process of

predicting FN. 49

D . Testing

1 . Tensile Testing at 4 K

The 4-K stress—strain curves, shown in figures 4 and 5, illustrate

several characteristic changes in the curves that take place as either the

test temperature is decreased or the austenitic alloying element concentrations are increased in metastable 18Cr-8Ni stainless steels. Considering

the test temperature effects first, figure 4 shows the occurrence of sudden

load drops or 'serrations that are associated with the 4-K stress-strain

curves. Serrated curves are common in Fe-18Cr-8Ni alloys at this temperature and have also been observed in other austenitic stainless steels,

brasses, bronzes, coppers, and aluminum alloys. The occurrence of these

serrations is a low- temperature phenomenon with the exception of materials

undergoing strain aging at ambient temperatures. 50

One explanation of the serration process is that a nucleation deformation event occurs as a result of local stress concentrations, thermal

fluctuations within the sample, or both. This event produces heat that is

not easily dissipated because of the low thermal conductivity of impure

metals at low temperatures. This results in a lowering of the local flow

stress, and deformation can continue at the applied stress level. This

produces more heat, and deformation continues until some undefined equilibrium is reached with respect to local heat generation, the absorption of

heat by the bulk specimen, and the applied stress. At this point a seem- ingly elastic reloading of the sample begins until another load drop takes

place. This adiabatic heating explanation is supported by the measured

mechanical energy dissipated in load drops, by correlations to thermal

energy required for the process, and by the ability to predict the temperature increase using specific heat, the temperature dependance of the flow

stress, and volume strained. 51

192
The character of the stress-strain fo^: these metastable 18Cr-8Ni

stainless steels is also dramatically altered by increasing the nitrogen

and manganese alloy contents. As the concentration for either of these

austenitizing elements is increased, the alloy becomes more stable against

stress- induced martensitic transformations. In figure 5, stress-strain

curves characteristic of an 18Cr-8Ni base material are shown. As the man- ganese content increased from 1 to 6 wt.%, the yield strength increased and

the work-hardening rate decreased.

2 . Charpy V-notch Testing

The Charpy V-notch impact test is commonly used to measure the ability

of a material to absorb energy in the presence of a flaw. The absorbed

energy is the sum of the crack initiation and propagation energies. It is

controlled by strength and ductility. The strength of the material regulates the force required to deform the specimen during testing. The

ductility of the material determines the distance through which the force

acts. The absorbed energy (AE) is a product of force times distance and

represents toughness in a manner similar to the area under a stress—strain

curve. Lateral expansion (LE) , on the other hand, is a direct measure of

notch ductility. So, since the area under the stress—strain curve may not

change substantially for materials of varying strengths, a loss in notch

ductility measured directly by a loss in lateral expansion may not result in

a measurable loss in the absorbed energy. For example, in a higher strength

material, an increased absorbed energy value is required to obtain the same

lateral expansion. This fact is illustrated by figure 6. 52 By replotting

this data as LE versus strength it is found that LE can decrease linearly as strength increases at a constant AE. Thus, lateral expansion and absorbed

energy measurements from Charpy V-notch tests may imply seemingly contradictory results for materials of different strengths.

II. EXPERIMENTAL PROCEDURES

A. Materials

Two SMA weld metal test matrices were evaluated in this study. The

first test matrix, referred to as the manganese—nitrogen matrix, was used

for mechanical testing. The second test matrix was used to evaluate the

effects of compositional components on the weld metal ferrite content.

The test welds for the manganese—nitrogen test matrix were produced

using a series of 3.2-mm (0.13-in) diameter electrodes that had core wires

from a single heat of type-308 stainless steel. Alloy additions were made

to the SMA electrode coating. The chemical compositions of the individual

weld deposits tested in this study are given in table 2. These compositions

represent a test matrix, based on 18Cr-8Ni stainless steels, in which man- ganese (1.5 to 10 wt.%) and nitrogen (0.04 to 0.26 wt.%) were the principal

193
variants of interest. Several molybdenum (2 and 4 wt.%) and boron (0.006

and 0.015 wt.%) additions were also made to the matrix. The chromium

content of the matrix was varied (15 to 21 wt.%) to compensate for the

manganese, molybdenum, and nitrogen additions so that welds with a nearly

constant ferrite content were produced.

The weld metal carbon and nitrogen contents were determined with individual element analyses. The phosphorus and sulfur contents were determined

by optical emission spectrophotometry. The remaining element contents were

determined by using conventional analytical techniques. The ferrite con- tents of weld pads and the welds were measured magnetically. The FN mea- surement device was calibrated with the ANSI/AWS standard A4. 2-74. 53

The second test matrix consisted of 929 SMA stainless steel weld com- positions (Appendix 1). Most of these data were supplied by electrode manu- facturers in cooperation with the Welding Research Council. The remainder

of the data was collected from the literature. The nitrogen—manganese test

matrix is included in these data. The FN values used for all welds in this

test matrix were those measured magnetically. The FN values ranged from

approximately zero to the very high values associated with duplex stainless

steels. The welds ranged in composition from commercial types 308, 316,

and 310 and duplex stainless steel alloys to experimental compositions.

The compositional ranges for the various elements considered in the study

are shown in table 3.

The data were divided into four groupings for analysis. Data having

FN values of less than one were not used during any of the regressions to

eliminate the uncertainties associated with very low FN measurements. The

set of data with this restriction was defined as group 1. In addition to

this overall limitation, further divisions of the data were made. Data

group 2 was limited to FN values between 1 and 30 and to compositions for

which molybdenum content did not exceed 2 wt.% when the chromium content was greater than 20 wt.%. This data group excluded experimental alloys of high

combined molybdenum—chromium contents and the duplex stainless steel com- positions. Several data groups were used to evaluate more specific types of

compositional ranges. Data group 3 incorporated the restrictions of group 2 and limited the nitrogen and manganese contents to 0.1 and 3.0 wt.%, respectively. These data groups were used to compare accuracies between FN equations for welds whose compositions were close to those of commercial weld

material. Types 308, 309, and 316 stainless steel weld compositions were

evaluated separately. Data group 4 was limited to welds having primary

ferritic solidification modes, as defined by Suutala. 49

B . Welding

Mild steel plate cut into 13-mm-thick by 305-mm-long coupons and beveled for a single vee weld joint configuration was used as a base material

To overcome the effect of dilution of the weld deposits by the base metal,

the exposed faces and backing strip were buttered with two layers of weld

194
metal prior to beginning the test weld. This procedure conforms to the

AWS A5.4-78 specification. 54 The same welder, power supply, and welding

parameters were used to minimize weld variations. Welding was done at

110 A and 22 V using an identical bead sequence. The heat input was approximately 0.6 kJ/mm, and an interpass temperature of approximately 93°C was maintained.

C . Mechanical Test Specimens

Two all-weld-metal 6 -mm- diameter 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.

Five Charpy V-notch specimens (ASTM E23 type A) were removed from each

weld. 55 Their notches were oriented perpendicular to the plate surface

along the weld center line and parallel to the direction of welding.

D . Tensile Testing

The testing apparatus used for the 4-K tensile testing is shown in figures 7 and 8. The upper specimen grip load line connects to a fixed upper

crosshead that contains a built-in load cell. The lower specimen grip load

line is anchored to the bottom of the compression tube, and as the crosshead

moves downward, the specimen is put in tension. A conventional stainless

steel Dewar, like the compression tube, is fixed to the moving crosshead and

is used as the liquid helium reservoir. Liquid helium is introduced to the

reservoir via an inlet passing through the crosshead and measured by a liquid helium superconducting liquid level indicator. The specimens are kept

submerged in liquid helium throughout the test.

The extension of the specimen within its gauge length is measured with

three cantilever beam extensometers seated on a retaining collar and positioned at 120 degrees with respect to each another. The retaining collars

are fixed to the specimen by tungsten—carbide- tipped set screws. A 24. 5 -mm

long fixture is used during the mounting to help maintain consistent gauge

lengths (figure 9) . The extensometers are accurate to approximately ten

percent strain. They are calibrated individually prior to testing at 76 K

with a vertically mounted dial micrometer. The extensometers are wired in

parallel so that an average of their outputs is recorded during the test.

A strain rate of 9.5 x 10" 5 s' 1 and a gauge length of 24.5 mm was used for

all tensile tests.

Two strip chart recorders were used. One recorded a load—time record

of the test and one recorded load—displacement up to approximately ten

percent strain. After ten percent strain this recorder was also used to

record load—time data.

195
E

.

Charpy V-notch Testing

Charpy V-notch testing was done at 76 K following the ASTM E23 test

standard method. A 163-J (120-ft-lbf) capacity unit was used. The impact

velocity was calculated to be approximately 4.8 m-s -1.

F. Linear Analysis

Stepwise linear analysis was done on a mainframe computer. Quantities

allowed as choices during the regression can generally be categorized into

three groupings: (1) simple compositional variables such as carbon (C)

,

nitrogen (N) , chromium (Cr) , and molybdenum (Mo) , (2) interaction variables

such as Mn x N, Cr x Mo, and Mo 2

, and (3) equation variables. The equation

variable grouping in this study was used to submit existing FN predictive

equations for evaluation. These equations were "seen" as a single variable

during the regression. The three equations included within this category

that are of the most interest here were the "DeLong-type FN equations" (FNCD

and FNCDK) and the FN predictive equation proposed by Hull (FNCH)

.

3

8

’ 3

9

> 4

4

The notations for these equations are defined in table 4.

The program repeatedly evaluated the variable list and added elements

and interactive terms to a predictive equation in the order of their statistical significance. In step one, the regression program selected the most

highly correlated variable for entry into the model. A coefficient and

intercept was then assigned. In step two the process was repeated, and a coefficient was assigned to the variable selected in this step. The coefficient chosen for the variable entered in step one was adjusted at this point

and a new intercept value was assigned to the expression. This process was repeated until none of the remaining variables could meet the F- statistic

selection criterion. For this study, an F value of 4.0 was chosen as the

critical value. The F value is defined as the regression coefficient

divided by the standard error squared. This value can also be expressed

roughly as a 95 percent confidence level or a ratio of the best parameter

estimate to a standard error of 4.0. The program also defined a multiple

correlation factor squared (R2

) and a standard error of estimate value to

each of the equations.

G

.

Microscopy

1 . Light Microscopy

For making general microstructural observations, samples were taken

from test specimens and prepared by standard metallographic techniques. A

mixed acid etch containing 55 percent hydrochloric, 22 percent nitric, and

22 percent acetic acids was used as a general etch. To highlight only the

delta ferrite phase, samples were electrolytically etched in a potassium

hydroxide solution.

196
2 . Scanning Electron Microscopy

After ultrasonic cleaning in methanol, fracture surfaces of broken

specimens were characterized using a scanning electron microscope (SEM)

Polished surfaces from the specimens were used for inclusion counts.

H. X-rav

Weld samples 6 and 13 were step scanned at intervals of 0.02 degrees,

29, for 25 s per step. A tube voltage of 44 kV at 40 mA was used. The

radiation was Cu Kp (ACu/3 = 1.39225 A). The specimen surfaces were prepared

by grinding followed by electropolishing.

III. RESULTS

A. Microscopy

1 . Light Microscopy

During light microscope examinations of mechanically polished weld

metal samples, all welds were found to have large numbers of inclusions.

In figure 10, a micrograph of alloy 1 shows the typical appearance of the

unetched weld deposits.

The ferrite morphologies observed after etching the welds were predominately a result of primary ferritic solidification. In figures 11 and 12,

micrographs of alloy 6 show the ferrite structures commonly observed in the

welds. The two principal ferrite morphologies found were vermicular and

lacy. Weld metal having more complex ferrite morphologies and growth orientations, however, was observed in areas located between welding passes. In

figure 13, an interpass region from alloy 22 is shown. At higher magnification, this high nitrogen—manganese sample was found to have some areas

of weld metal that solidified by primary austenitic solidification. In

general, however, most of the welds did not appear to change or have mixed

solidification modes, even in these complex interpass areas.

2 . Scanning Electron Microscope Inclusion Counts

Inclusion distributions were plotted for several representative welds

from data collected at high magnification on the SEM. In figure 14, the

inclusion distributions for the welds are represented by an envelope indicating the variations between samples. Alloy 6, the 3.84 wt.% molybdenum

alloy, was found to have a large number of small-diameter inclusions.

Micrographs of alloys 15 and 6 taken at the magnification used during the

counts are shown in figures 15 and 16. The average inclusion diameters

were approximately 3.3 /im, however, the variation in size was substantial.

Almost all the inclusions were spherical in shape.

197
B . Strength

Stress—strain curves for five alloys that characterize the observed

changes in the curves as a function of alloy content are shown in figure 17.

As alloy content increased, the yield strength increased, the work-hardening

rate decreased, and the distinctive three -stage curve was lost. These

curves are shown without serrations for simplicity. In figure 18, a representative serrated curve is shown.

The specific yield and tensile strengths determined for the various

welds at 4 K are given in table 5 and plotted versus nitrogen content in

figure 19. The tensile strength remained nearly constant at an average of

1400 MPa for nitrogen contents up to 0.16 wt.% and increased slightly to an average of 1500 MPa for the three alloys having about 0.26 wt.% nitrogen.

The differences in tensile strength among welds were not statistically re- lated to alloy content. The yield strength, however, was clearly related

to the alloy content of the welds. The dominant strengthener was nitrogen.

The equation expressing the yield strength solely as a function of nitrogen

content was

Weld Yield Strength (MPa at 4 K) = 400 + 3600(wt.% N) (6)

The F value (a test of the significance of the model) was 265. The coefficient of determination (R2

, a measure of the quality of the regression) was

0.92. The standard error of the prediction (STD) was 83 MPa.

Increases in yield strength due to manganese additions are shown in

figure 20. At the 0.05, 0.10, and the 0.15 wt.% nitrogen concentrations, an approximate 15 MPa increase in yield strength per wt.% increase in manganese

content was found. This coefficient represents an average of the three individual slopes and assumes no nitrogen—manganese interaction.

A graphical separation of the effects of manganese and nitrogen from

factors such as molybdenum content, slight variations in FN, and other

compositional elements becomes difficult at this point. Using linear

stepwise regression, the equation best able to predict the yield strength

was

:

Weld Yield Strength (MPa at 4 K) = 334 + 3407(wt.% N) + 15(wt.% Mn) -l- 46(wt.% Mo). (7)

2 This equation had an F value of 591, an R of 0.96, and a standard error of

estimate of 57 MPa.

C. Toughness

1 . Absorbed Energy

The overall CVN absorbed energy at 76 K decreased slightly as the

nitrogen content was increased from approximately 0.04 to 0.16 wt.% (fig- ure 21). The average absorbed energy value within this region was 40 J.

198
Varying effects of nitrogen were found when alloys of otherwise similar

composition and FN were compared at 0.04, 0.10, and 0.15 wt.% nitrogen

levels. For example, at 1.57 wt.% manganese, alloys 1 and 7 had respective

CVN absorbed energies of 46 and 39 J. This implies a loss of approximately

12 J per 0.1 wt.% nitrogen. At 3 wt.% manganese, alloys 18, 17, and 11

showed little reduction in absorbed energy as a function of nitrogen content. At 6 wt.% manganese (alloys 4, 9, and 12), the loss in absorbed

energy per 0.1 wt.% nitrogen was approximately 5 J.

Nitrogen contents above 0.16 wt.% caused a substantial decrease in

absorbed energy. At 0.25 wt.% nitrogen, the absorbed energy was by reduced

by approximately 50 percent with reference to the 40 J average at lower

nitrogen contents.

Although some indications of reductions in absorbed energy due to

solid- solution alloying additions and FN were present, with the exception

of the 3.84 wt.% molybdenum sample, no clear trends were resolvable within

the data. Comparing the high molybdenum sample (alloy 6) to the 1.66 wt.%

molybdenum sample of similar manganese, nitrogen, and ferrite content (alloy

13)

a decrease in absorbed energy of 17 J was found.

2 . Lateral Expansion

The lateral expansion decreased linearly as a function of nitrogen

content, and to some extent as a function of manganese and molybdenum contents. Slight variations in the ferrite content of the samples could not

be correlated to changes in the lateral expansion measurements.

As shown in figure 22, a decrease in lateral expansion of approximately

86 percent occurred when the nitrogen content was increased from 0.03 to

0.26 wt.%. Allowing nitrogen as the only compositional element to be chosen

in the stepwise linear regression, the resulting equation was:

LE (mm) = 0.686 - 2.11(wt.% N) (8)

2 This equation had an F value of 91, an R value of 0.79, and a standard

error of 0.081 mm.

It is apparent from figure 22, however, that manganese also contributes

to the loss in lateral expansion at a given nitrogen content. The lateral

expansion of alloy 12, for example, is 0.1 mm less than that of alloy 11.

Both alloys have 19 wt.% chromium, 0.15 wt.% nitrogen, and 1.5 FN. The

loss in lateral expansion is a result of the 3.4 wt.% manganese increase in

alloy 12.

The alloys with molybdenum additions generally fit the observed trends;

however, the lateral expansion values of alloys with combined high nitrogen,

manganese, and molybdenum content (particularly alloys 13 and 16) were low.

Allowing all compositional elements and FN as acceptable variables, an equation expressing the effect of molybdenum, manganese, and nitrogen on lateral

expansion was found.

199
( 9 ) LE (nun) = 0.787 - 1.81(wt.% N) - 0.031(wt.% Mo) — 0 . 022 (wt . % Mn)

.

This equation had an F value of 280, an R

2 value of 0.92, and a standard

error of 0.051 nun.

Since the lateral expansion was found to decrease linearly with increasing alloy content, it was inversely proportional to the yield strength.

The strength—toughness relationship is shown in figure 23. The equation

found to express lateral expansion in terms of the 4-K yield strength was:

LE (mm at 76 K) = 0.95 — 6 x 10' 4 (Yield Strength, MPa) (10)

This equation had an F value of 344 and an R

2 value of 0.94.

D

.

Fractography

Macroscopic examinations of fracture surfaces showed texture differences on Charpy V-notch specimens of varying nitrogen and ferrite contents.

In general, the higher nitrogen content samples had coarser, less planar

fracture surfaces. At nitrogen contents greater than 0.15 wt.%, the fracture surfaces with the most texture also had the lowest FN (figures 24a

and b) . Shears lips tended to be present only on a Charpy V-notch specimen

having nitrogen contents of greater than 0.15 wt.%. Examination of the

tensile specimens at low magnification showed little or no correlation

between fracture surface appearance and alloy content. All samples had

fine, fibrous, planar fracture surfaces. Little necking took place in the

tensile specimens. A narrow shear- lip region was present on many of the

samples, but no relationship between shear- lip formation and nitrogen

content was observed.

Examinations at high magnifications (SEM) of the Charpy V-notch and

tensile specimen fracture surfaces showed the failure mechanism in all

samples to be predominantly void nucleation and growth. Typical ductiledimple tear-ridge networks were observed on all fracture surfaces. The

tear-ridge patterns were similar to the two major orientations of ferrite

present in the welds (figures 25 and 26). The results of dimple counts

taken at 3000x on fracture surfaces showed some variation in dimple densities; however, compositional variations did not alter the inclusion density

or size noticeably (table 6). Ductile dimple size was very similar between

the various welds and specimen types. Most surfaces had a fine dispersion

of small round inclusions and occasional large inclusions. In figure 27

is shown a typical dimple whose size is typical for the welds tested in this

s tudy

.

The 3.8 wt.% molybdenum sample had a greater number of large inclusions

on its Charpy V-notch fracture surfaces. These inclusions had a glassy,

200
brittle appearance and were found to be rich in calcium, manganese, silicon,

and titanium during energy dispersive analysis in the SEM (figure 28).

The ductile dimples of tensile specimens had the same general size and

distribution as those observed on the Charpy V-notch specimen surfaces

(figure 29). There were, however, several types of characteristic fracture

features seen only on Charpy V-notch fracture surfaces (figures 30 and 31).

Both types of specimens had isolated areas where fractures of a more brittle

nature occurred. Areas clearly fractured by cleavage in both sample types

were predominantly ferrite (figures 32 and 33).

E . X-ray Analysis

The types of phases present in the two welds analyzed (alloys 6 and 12)

were identical. The relative quantities of the phases differed slightly

(figure 34). The principal phase present was austenite.

F . Ferrite -Number Prediction

Four FN predictive equations were developed in this study. These equations (FN1, FN2 , FN3 , and FN4) are defined in table 7, which also gives the

F, R

2

, and standard error of estimate values.

Comparing the fits of FN1 , FN3 , and FN4 equations to those of the

DeLong-type equations (FNCD, FNCDK) and to the Hull equation (FNCH) , the

accuracy of FN prediction was improved for the welds of data group 2. The

FN3 equation had an R

2 value of 0.73 and a standard error of estimate of

1.7 FN. This -equation more accurately predicted FN than any other equation

developed or considered over the wide range of alloy contents present in the

group 2 weld data.

Considering only welds having compositions close to those of commercial

alloys (group 3), the predictive accuracy of the DeLong-type equation improved. The FN1 equation, however, was a slightly better predictor of FN

than the DeLong-type equation for this data group.

The FN2 equation was developed from weld data sorted to include only

those welds that had primary solidification modes. It was found that the

coefficients within this equation changed with respect to those of the FN1

equation. Most notably, the coefficient for nitrogen increased, and the

coefficient for carbon decreased.

201
IV. DISCUSSION

A. Strength

1 . Stress—Strain Curves

The changing character of the stress-strain curves plotted in figure 17

indicates that the 4-K stability of the welds with respect to martensitic

transformation is increasing with nitrogen and manganese alloy content. The

three-stage curve representing alloy 1 resembles the curve typifying 18Cr8Ni stainless steel base material (figure 5). At low nitrogen and manganese

contents (alloy 1), deformation occurring at low strain required only slight

increases in stress to continue. It has been proposed that deformation at

nearly constant stress is possible in stainless steel base materials due to

the formation of bcc martensite at slip-band intersections. 50 The laths act

as "dislocation windows" between slip planes and enable the deformation pro- cess to continue without work hardening. The increased work-hardening rates

at higher strains are associated with hep martensitic formation. Although

the curves are only approximations and the work hardening rates are somewhat

uncertain, it is clear that increasing the manganese content, nitrogen con- tent, or both alters the deformation process in the weld metal. The absence

of the "easy glide" region for these alloys would be attributed to the suppression of bcc martensite formation at high interstitial and solid- solution

alloy content.

In alloy 1, following the "easy glide" region, the substantially higher

rate of work hardening is associated with a higher percentage of austenite

to martensite -transformation. In the more stable alloys (10, 12, and 19),

however, there is still appreciable martensitic transformation occurring.

Alloy 22 (0.259 wt.% nitrogen, 9.7 wt.% manganese) is on the verge of being

completely stable. In agreement with the finding on base metals, the work

hardening rate decreases with increasing austenite stability.

2 . Yield and Tensile Strength

The 4-K yield strength as a function of alloy content was expressed

reasonably well by equations 6 and 7. Equation 7, by also including man- ganese and molybdenum factors, was a better fit to the data.

Weld 4-K Yield Strength (MPa) = 334 + 3407 (wt.% N) + 15 (wt . % Mn) + 46 (wt . % Mo). (7)

Large intercept and error values, however, point out that slight variations

in ferrite, chromium, and carbon contents are influencing the strength and

are beyond the resolution of this study. Nonetheless, it is clear that

interstitial nitrogen is the dominant strengthener and that molybdenum has

slightly more strengthening power than manganese. The relative magnitude of

the nitrogen coefficient proposed in equation 7 is supported by studies on 18Cr—8Ni base material and on 19Cr—14Ni SMA weld metals. 56

202
Base Material 4-K Yield Strength (MPa) = 350 + 3400(wt.% N)

( 20 )

19Cr—14Ni SMA Weld Yield Strength (MPa) = 145 + 3441 (wt.% N) + 25 (wt . % Mo) + 15 (wt . % Mn) + 8122(wt.% C) . (21)

where F = 2416, R

2 = 0.997, and STD = 12 MPa for equation 21.

The molybdenum coefficient of 46 proposed by this study appears too

high when compared to that of equation 21. Because of the limited number

of molybdenum alloys tested, it has the most uncertainty associated with

it. There is some indication, however, that the strengthening potential of

molybdenum decreases as nickel contents increase at 4 K. 57 This would support the finding of a larger coefficient for molybdenum at 9 wt.% nickel

than at 14 wt.% nickel.

Other interactions between compositional elements that alter the

strengthening character for one alloying element as a function of overall

alloy content may also exist. For example, in figure 20 the slopes associated with the strengthening power of manganese appear to decrease as the

nitrogen is increased from 0.05 to 0.10 to 0.15 wt.% levels. This implies

that the strengthening coefficient of 15 attributed to manganese in equation 7 represents the average effect of manganese throughout the nitrogen

range of the test matrix. The actual coefficients may vary depending on

nitrogen content. Possibly a more accurate expression of yield strength as

a function of composition would need to have manganese—nitrogen, molybdenumnickel, and additional cross terms associated with it. For a nitrogen SMA

weld having approximately 3 wt.% manganese and a low molybdenum content,

however, it is felt that equation 7 would predict the 4-K yield strength

accurately

The results showed that variations in the 4-K tensile strength are not statistically related to alloy contents. For example, alloy 1 (1.57

wt.% Mn, 0.047 wt.% N) and alloy 22 (9.72 wt.% Mn, 0.259 wt.% N) had respective tensile strengths of 1586 and 1498 MPa. Inhomogeneities within the

weld structures, as shown in figures 32 and 33, and variations in ferrite

contents are suspected to be masking any effect that manganese, molybdenum,

or nitrogen have on the 4-K tensile strength. Recent studies conducted with

fully austenitic weld metals show that tensile strength increases slightly

with increasing nitrogen content. 66

In summation, as the interstitial and solid- solution alloy content increased in the welds, the yield strength increased linearly and the tensile

strength remained relatively constant. The fact that the strengthening co- efficient defined for nitrogen in this study agrees well with those found in

base metals and other stainless steel weld compositions implies that nitrogen additions act principally in an interstitial strengthening capacity. If

appreciable amounts of nitrides or other possible forms for nitrogen in the

microstructure were being formed, a difference in nitrogen strengthening as

a function of overall weld composition might be expected. The relatively

small influence of alloy additions on the fracture stress and strain points

out that the failure mode (void nucleation and growth) was not altered by

increased alloy contents.

203
B . Toughness

There are two types of questions raised by the Charpy V-notch test

results in this study. Questions in the first category are related to the

differing results, concerning the effect of nitrogen on toughness, found

from absorbed energy versus lateral expansion measurements. Should the

relationship between these measurements of toughness be linear? If they

should, which alloy compositions deviate from the expected trend and why?

Questions within the second category are related to the more specific

effects of compositional on the toughness. In particular, is the dramatic

decrease in absorbed energy found at high molybdenum contents real?

1 . Differences between Measures of Toughness

Addressing the question of linearity between lateral expansion and

absorbed energy first, the data from this study are compared with data from

a study27 in which linear relationships were proposed for SMA type -308 and

type-316 stainless steel welds at 76 K. In figures 35 and 36 the data sets

are shown plotted as absorbed energy versus lateral expansion. Lines representing the expected linear relationships are also shown in the figures. It

should be noted that data plotted here from the Szumachowski—Reid studies

were selected to correspond to the nickel and ferrite contents of the welds

tested in this study (Ni 9 wt.%, FN < 6). The compositions of these welds

are given in table 8. From the type- 308 and -316 stainless steel trend

lines alone, one would suspect that the welds which deviate from the ex- pected trends would be those making up the knee of the curve. In both data

sets these welds tend to be of medium nitrogen content (0.08 to 0.15 wt.%)

and have substantial manganese content, molybdenum content, or both.

Because both data sets show tendencies toward the same type of deviation from the proposed linear relationship between lateral expansion and

absorbed energy, it is concluded that linearity between these two measures

of toughness from the Charpy V-notch test should not be expected for highly

alloyed welds of varying composition.

It is felt that the differences between these two measures of toughness can be related, at least in part, to the varying strengths of the

alloys tested. In steels for example, as discussed in the introduction,

material of increasing strengths have similar absorbed energies but decreasing lateral expansion values. This is because absorbed energy is defined by

the product of force times distance, whereas lateral expansion is a direct

measure of ductility. Although ultimate strength is traditionally used in

these arguments, a similar trend might be expected for increasing flow

strengths [ (ay + /2].

Because these two measures of toughness differ for the alloys, both

types of results deserve equal attention. Relationships between fracture

toughness and impact test results are not known at this time. So even

though lateral expansion is a more direct measure of ductility, absorbed

energies may relate additional information concerning crack growth.

204
2 . Effects of Alloying on Toughness

The lateral expansion results (figure 22) clearly show that as nitrogen

contents increased, notch ductility decreased. At nitrogen levels greater

than 0.16 wt.%, the absorbed energy results also show the adverse effect of

nitrogen on toughness (figure 21). At lower nitrogen contents, the absorbed

energy results are somewhat more difficult to interpret. Considering alloys

1 and 7, for example, which have approximately 1.5 wt.% manganese, a 12 -J

decrease in absorbed energy per 0.1 wt.% nitrogen was found. This agrees

well with the data of Szumachowski and Reid. 4 Trends indicated by alloys

11, 17, and 18 or alloys 4, 9, and 12, however, imply that the detrimental

effect of nitrogen on toughness for higher manganese alloy contents was

reduced. Similar trends were indicated by the lateral expansion results.

The lateral expansion was shown to decrease as manganese contents increased

at a given nitrogen level, but at approximately 1 wt.% nitrogen, the increase in manganese content from 1.5 to 3 wt.% did not result in additional

notch ductility losses. Increasing manganese contents further at this

nitrogen level to increase strength resulted in nearly equivalent losses in

toughness on a percentage basis.

In a more general manner, equation 9 can be used to compare the

relative effects that nitrogen, manganese, and molybdenum have on notch

ductility

LE at 76 K (mm) = 0.78 — 1.8 (wt.% N) — 0.03 (wt.% Mo) — 0.02 (wt.% Mn) . (9)

For example, losses in lateral expansion of 0.15 and 0.36 mm, respectively,,

would be predicted if the lowest manganese and nitrogen concentrations in

the test matrix were increased by a factor of 5. This translates to

approximate losses in lateral expansion of 33 percent for the 7.5 wt.%

manganese alloy and 51% for the 0.2 wt.% nitrogen alloy. The significant

predicted loss in lateral expansion due to manganese again points out that

optimization of solid- solution alloy contents with respect to nitrogen con- tent deserves future consideration. Overall, the equation expresses that

nitrogen has the most detrimental effect on lateral expansion, followed by

molybdenum, and then manganese. As might be expected, nitrogen effected the

greatest increase in strength, followed by molybdenum, and then manganese

(equation 7)

Compositional coefficients in equation 9, however, only represent the

averaged effects of these elements on toughness. The molybdenum coefficient

in particular is somewhat uncertain, since at 3.84 wt.% a sudden decrease in

toughness was found and a limited number of molybdenum alloys were tested.

This may skew the results and misrepresent the effect of molybdenum on the

notch ductility.

The reason for the significantly lower toughness of the 3.84 wt.%

molybdenum alloy was not clearly determined. The adverse effect of high

molybdenum content, however, was confirmed by testing a second weld whose

205
composition was similar to that of alloy 6. In an effort to understand the

degradation associated with these higher molybdenum alloys, segregation,

microstructural , and fractographic studies were conducted.

Welds 6 and 13 had very similar compositions with the exception of

their chromium and molybdenum contents. Weld 6, a 15 wt.% chromium and 3.84

wt.% molybdenum alloy, had an absorbed energy of 16 J. Weld 13, a 18 wt.%

chromium and 1.66 wt.% molybdenum alloy had an absorbed energy of 33 J. Uhe

loss in absorbed energy of approximately 50 percent is attributed to the

increase in molybdenum from 1.66 to 3.84 wt.%. Micros trueturally, the welds

were very similar. During examinations with a light microscope, both welds

were found to have roughly equivalent ferrite contents and primary ferritic

solidification morphologies. No evidence of intermetallic phases was found.

X-ray analyses of chemically polished weld surfaces were inconclusive.

Inclusion counts on polished surfaces by SEM did show that the high molybdenum weld had a slightly different inclusion size distribution that of

other welds analyzed from the test matrix. It had a larger number of smalldiameter inclusions and a higher volume fraction of inclusions. Examination

of tensile and Charpy V-notch fracture surfaces showed that all welds failed

in a ductile manner by microvoid coalescence. The 3.84 wt.% molybdenum

samples, however, had some very large inclusions on their fracture surfaces

(figure 28) . These inclusions had: diameters ranging from 20 to 100 /im;

compositions rich in calcium, manganese, silicon and titanium; and glassy,

conchoidal fractures associated with them. It might be speculated that the

difference in the 3.84-wt.% molybdenum- sample inclusion- size distributions

may be related to changes in the weld pool viscosity. Molybdenum has a high

melting point (2610°C) . This may increase the solidus temperature of the

system and result in an increased weld-pool viscosity. In this case, there

would be less opportunity for small inclusions to coalesce and a larger

chance for large inclusions to be trapped. Whether the presence of more

small- and large - diameter inclusions in alloy 6 was solely responsible for

lowering the toughness in these samples is not known.

In these final paragraphs of this section, the strength—toughness

relationship determined for the weld metals is summarized. The inverse

relationship found between lateral expansion and yield strength at 4 K

(see figure 23) was best expressed by equation 10:

LE (mm at 76 K) = 0.95 - 6 x 10' 4 (Yield Strength, MPa) (10)

Perhaps most importantly, this equation shows that notch ductility is

principally a function of strength. Interstitial strengthening elements,

such as nitrogen, greatly increase strength and reduce notch ductility proportionately. Solid-solution-strengthening elements, such as molybdenum

and manganese, strengthen the weld metal to a much lesser extent. Therefore, they are less detrimental to the notch ductility. Other equations

developed in the study (7 and 9) serve to help quantify the relative effect

of the alloying elements on strength and toughness. Previous discussions

have pointed out that optimizing manganese contents at a given nitrogen

level may slightly increase toughness and that molybdenum contents near

4 wt.% may severely decrease toughness in type-18-8 stainless steels weld

206
deposits. Overall, however the alloys tested in this study deviate little

from the linear relationship between lateral expansion and yield strength

defined by equation 10. To significantly improve this relationship, factors

other than manganese, nitrogen, and molybdenum alloying effects must be

considered.

Comparing the toughness values of the welds tested in this study with

those of the typical type-316LN stainless steel base material shows that the

strength-toughness relationship determined for these welds is too low. At

a 1000 -MPa yield strength the base material would be expected to have Charpy

V-notch absorbed energy and lateral expansion values of approximately 90 J

and 1.0 mm, respectively. The toughest weld (alloy 12) tested in this

study, with a yield strength above 1000 MPa, had an absorbed energy value

of 37 J and a lateral expansion value of 0.32.

Areas of study that may result in further toughness increases in SMA

type-18-8 stainless steel welds include methods to lower the weld metal

inclusion contents. Like inclusions, the ferrite content of the weld metal

also influences toughness. Although void nucleation did not occur at the

ferrite—austenite interface in this study, figures 25 and 26 imply that the

ferrite may play an important role during void growth owing to its influence

on plastic flow in the matrix.

The importance of controlling and predicting the FN for weld metal

alloys is demonstrated by (1) ferrite affected fracture surface appearances

for the low FN weld matrix test in this study, and (2) increased ferrite

contents reduced toughness in previous studies.

C . Ferrite-Number Prediction

The four FN predictive equations developed in this study (FNl , FN2

FN3 , FN4) are discussed in this section. The accuracy of predicting FN

using the FNl linear equation form is compared in detail with the accuracy

of FN prediction using previously developed equations. The FN2, FN3 , and

FN4 equations are discussed to highlight the further improvements in FN

prediction that are possible by the use of nonlinear equation forms and the

consideration of weld solidification mode prior to analysis. Before beginning the discussion, however, factors complicating the comparisons between

the predictive equations need to be identified.

1 . Fairness of Comparisons between Equations

Several factors complicate any direct comparisons between equations

developed in this study to previously existing equations. As shown in the

results (table 7), equations representing the DeLong diagram (FNCD, FNCDK)

and the equation incorporating the Hull coefficients (FNCH) are treated as a single variable by the regression program. They are fitted with slopes and

intercepts that improve their fit to the data. Therefore, the statistics

that characterize the significance of these equations are overestimates.

207
However, this error is partially balanced by the fact that mathematical

expressions representing each diagram differ slightly from the actual diagrams. For example, these equations do not duplicate the slightly decreasing slopes and spacing between increasing constant FN lines of the original

DeLong diagram.

Since the error between the DeLong- type equations and the DeLong diagram is small at FN values less than 10 and most of the data in this study

have FN values close to or less than 10, the error associated with this

factor is believed to be small. For this reason, figures representing the

DeLong- type equation are used in this discussion. It is believed that they

enable the most direct comparisons between equations to be made. These

figures represent the equations without the added complication of the re- gression program making adjustments to their respective fits.

2 . Ferrite-Number Equations

a) Linear Equations

The DeLong nickel and chromium equivalents, developed for low nitrogenmanganese alloys, increased the error associated with FN prediction in more

highly alloyed weld metals. The DeLong manganese term (0.5 * Mn) was a major contributor to the errors in FN prediction. In figure 37, a large

number of weld compositions having AFN errors greater than ±5 are positioned

at the upper left portion of the data. After substitution of the 0.34 con- stant for the DeLong manganese term, most of the error associated with this

region is eliminated (figure 38)

.

By representing the austenitizing power of manganese as a constant, the

tendency to underestimate the predicted FN is reduced for high manganese

alloys. Still a substantial number of high AFN error exists in figure 38.

By using equation FN1 , further reductions in the remaining high AFN errors

were made, as shown in figure 39. Approximately 75 percent of the welds had

predicted FN values within 2 AFN when the FN1 equation was used. Using

either the FNCD or FNCDK expressions, only around 50% of the data fell within this error range (figure 40). The most significant improvement, resulting from the replacement of the DeLong manganese term by the 0.34 constant

in the FNCDK expression, was the reduction of the greater than 5 AFN error

from 20% to almost 10%. Many of the welds that had been in this high AFN

error group were moved into the 3 < AFN < 5 error group. Since neither the

FNCDK nor the FN1 equations have manganese terms, further reductions in AFN

error found when the FN1 equation was used are attributed to differences

between other terms in the two equations.

Comparison of individual compositional coefficients between equations

is difficult. The magnitudes of the coefficients within a given equation

are relative to one another and the intercept values between equations vary.

In table 9, the unfactored coefficients for the various equations are given,

however, and interestingly the coefficients for nickel are quite similar.

The chromium coefficient, or more importantly the difference between the

208
chromium and nickel coefficients, is the greatest for the DeLong-type

expressions. If the 3.5 chromium coefficient is incorrect, terms within

the nickel equivalent expression must have increased magnitudes to offset

the error. If the nickel coefficient is approximately correct, this leaves

the nitrogen, carbon, and manganese terms responsible for much of the

correction.

Examining the compositional coefficients within the various predictive

equations showed that many of the DeLong coefficients do not adequately

model compositional effects on FN. The 3.5 chromium coefficient, for

example, resulted in increasingly positive AFN deviation as the chromium

content was increased (figure 41). For chromium equivalent components,

positive AFN error implies that the coefficients are too high. The opposite

is true for nickel equivalent terms. The ferritizing power of chromium was modeled better by the FN1 equation in figure 42. In this equation the chromium coefficient is lower and there is less difference between its magnitude

and that of the nickel coefficient. These results imply that the FN predictive equations in which the respective ferritizing and austenitizing

powers of chromium and nickel are similar in magnitude can more accurately

predict FN.

The effects of molybdenum on FN, as modeled by the different predictive

equations, were examined in figures 43 through 46. In figures 43 and 44,

wt.% molybdenum is plotted versus AFN. No clear correlation between molybdenum and AFN deviations for either FNCD or FN1 was noted when the duplex

stainless steel compositions and other very high molybdenum—chromium con- tents were ignored (figure 45) . When the high molybdenum—chromium alloys

were included (figure 46) , the nonlinear relationships between AFN and

molybdenum content could not be modeled by a simple compositional term. It

appears that accurate prediction of FN in high chromium—molybdenum alloys

requires that nonlinear interaction terms be incorporated into the equation.

A final observation with respect to these figures: the DeLong-type expression causes the data to be offset to a relatively constant positive AFN

value. The FN1 equation positions the data more symmetrically about the

0 AFN line. This is evidently due to a difference in intercept values

between the two equations

The nitrogen coefficients in the nickel equivalent terms are evaluated

next. A trend toward increasingly negative AFN deviations was found when

the FNCDK expression was used to predict FN (figure 47) . The DeLong nitrogen coefficient of 75 is shown to be too high by this result. Figure 48

shows that the FN1 equation with its lower nitrogen coefficient is a better

model of the effect of nitrogen on FN. Again, additional figures to help

identify the high chromium—molybdenum (figure 49) and duplex stainless steel

data (figure 50) are presented. Probably the most noticeable difference in

the nitrogen coefficients of equation FN1 and the DeLong-type of equation is

that in the FN1 equation, nitrogen was modeled as being only half as powerful as an austenitizing agent as carbon, and DeLong' s coefficients attribute

equivalent austenitizing powers to the two elements.

209
The AFN error associated with the use of the (0.5 * Mn) DeLong manganese term is again shown to be significant in figure 51. Clearly, large AFN

errors result when this coefficient is used in the prediction of FN for

alloys having high manganese content. Although the absence of a manganese

term in FN1 greatly reduced predicted FN error (figure 52), the true effect

of manganese on FN has not been identified by the equation. The modeling of

the manganese effect by the Hull equation (FNCH) may come closer to expressing the true effect of manganese on FN. As manganese content increases, the

negative manganese- squared term reduces the effect on FN. The change in the

austenitizing power of manganese reflected in the FNCH expression is supported by solidification studies. 32 The correct modeling of higher order

terms, such as these and interactions between compositonal elements, is

difficult. This approach is discussed further when the nonlinear equation

forms developed in this study are presented.

The FN1 equation indicated that the DeLong coefficient for silicon is

too high and that the effect of titanium on FN may be substantial. For

better definition of the effects of these elements on FN, however, additional data would be required.

b) Special Cases

Equation FN1 predicted FN better when only welds of conventional alloy

(group 3) content were used to compare fits between it and the DeLong equation. The differences between the predictive abilities of the respective

equations, however, were greatly reduced when only low nitrogen and manganese contents were allowed. Two important points were brought out by this

comparison: (1) the DeLong diagram is an accurate method of predicting FN

for alloys within its suggested compositional range, and (2) that FN1 was developed over and extended compositional range has not lessened the ability

for this equation to accurately predict FN in the more commonly used stainless steel alloys.

In figures 53, 54, and 55, the relative accuracies of the FN predictive

equations are compared using type-308, -309, and -316 stainless steel welds.

For type-308 and -316 stainless steel weld deposits, both the DeLong- type

equation and the FN1 equation predicted FN values within 3 FN of the mea- sured values for over 95 percent of the welds considered. For type- 309

stainless steel alloys, however, the DeLong- type equation predicted FN

within the 3 FN range for only 53 percent of the welds. The FN1 equation

remained accurate within the 3 FN range for 95 percent of the type -309

compositions considered. So, the FN1 equation has, to some degree, increased the accuracy of FN prediction for the low alloy compositions.

Welds used to develop the FN1 equation were also separated on the basis

of solidification mode by using the Suutala equation. Welds identified as having primary ferritic solidification modes were then used as the data for

further linear regression analysis. The FN predictive equation determined

using this data (group 4) was FN2 . Comparing equation FN2 to FN1 : the

difference between the chromium and nickel coefficients decreased; the

210
relative austenitizing power attributed to nitrogen increased; and the aus- tenitizing power attributed to carbon decreased. Although the reported R

2 and F values are similar for the two equations (the lower number of cases

analyzed for FN2 tends to lower the F) , the decrease in the standard estimate of error and the differences between compositional coefficients are believed to be significant.

The question raised here is: can welds having different modes of solidification have their FN values predicted by the same equation? It seems

that FN cannot be singularly defined in terms of composition until this

ambiguity is dealt with. Primary ferrite and austenitic solidification

modes result in two physically distinct forms of ferrite.

c) Nonlinear Equations

An attempt was made to improve the fit for the extended data (group 1)

by allowing nonlinear terms. Equations FN3 and FN4 are discussed here to

help identify alternate FN predictive equation forms and possible interaction terms which may be necessary for better prediction of FN.

The FN3 equation was developed after variables expressing interactions

between compositional elements and self - interactions were permitted as choices during the regression using data group 1. This equation was the

most accurate predictor of FN for both the group 1 and the group 2 data.

Although the FN3 equation improves FN prediction, it is presented here

only to indicate the types of interaction terms that were consistently

deemed significant during the regressions. Principally, titanium—carbon

manganese—nitrogen, nickel—nitrogen, nickel—manganese , chromium—molybdenum

and nickel—chromium interactive variables were chosen during the regression.

Indications of the occurrence of interactions between compositional

elements in the literature supports the complex nonlinear form of equation

FN3. The work done by Strauss and Maurer 41 shows chromium to behave in a nonlinear manner with respect to nickel and ferrite formation. Previous

evaluation of the manganese—nitrogen test matrix used in this study by

Lake 42 suggested that including Mn — Mn2 — MnN + (MnN) 2 terms to represent

a nitrogen-manganese interaction best enabled FN to be predicted. In figure 56, a contour map of the nickel equivalent incorporating these terms

expresses that the austenitizing power of manganese changes at approximately

4 to 5 wt.% and that as nitrogen content increases, the effect of manganese is reduced. Solidification studies done by Suutala 40 support the

empirical results of Lake and other similar findings. 44 ’ 54

Inclusion of interaction terms such as these can clearly increase

the accuracy of FN prediction. The complexity of equation FN3 , however,

introduces uncertainty about which terms are truly expressing actual interactions. For example, if an interaction term is added to the equation

incorrectly, then other complex terms need to be added to correct for the error. For this reason, the primary element coefficients must first be

properly defined and possible interactions existing between primary elements

211
identified scientifically before the empirical development of more complex

FN equations is appropriate.

FN4 was developed using only primary compositional terms. This equation has the pseudo polar form of the Schaeffler equation. This form is

of interest because it can model the changing slopes and spacing of constant

FN lines found on DeLong- and Schaeffler- type diagrams. Using only simple

variables such as FN1 , it is suspected that this equation form can improve

the accuracy of FN prediction. Although the equation form was difficult

to refine using linear regression techniques, by assuming values for the

coefficients and constants included in the chromium and nickel equivalent

expressions, FN4 predicted FN better for data group 2 than FN1 . The

analysis indicated that particularly the nitrogen and titanium coefficients

required adjustment. However, nonlinear regression analysis can model the

coefficients better and develop a better predictive FN equation.

3 . Physical Significance

Clearly, it is more difficult to identify the nonlinear equation

forms and interaction terms that lead to improved FN prediction. Yet, one anticipates nonlinear behavior in complicated systems, so the eventual

development of these complex equations will not only improve the accuracy

of FN prediction, but will bring physical significance to these empirical

expressions

.

Ultimately, FN predictive equations should take on forms suggested by

the fundamental thermodynamic and kinetic considerations. A form suggested

by Liu et al. 39 to provide more fundamental guidelines for future predictive

equation development is:

FN = KNi Ni + KMn Mn + KCr Cr + % N + i (5) + KNi-Mn MnNi + KNi-Cr NiCr + Ki-j ij

where the coefficients (K^) include thermodynamic effects such as the con- figuration of atoms in the solution and parameters that compare bonding

energies between dissimilar and like atoms. In addition to defining the

equation coefficients, equation 5 also indicates that both linear and non- linear contributions are important.

Terms in equation 5, such as Ni and Cr, have long been recognized as important to austenite-ferrite stability. Nickel, for example, is an austenitic stabilizer and segregates preferentially to austenitic regions

during solidification. In austenitic stainless steel compositions, however,

alloy contents exceed those valid for dilute solutions, and therefore,

linear behavior should not be expected. Interactions between alloying elements in these complex solid solutions may occur and are indicated by terms

in equation 5 like Ni Mn. Proper incorporation of these interactive

terms into fundamental equation forms may help indicate the extent of clustering, compound formation, and other types of interactions existing among

atoms in these stainless steel compositions.

212
In this study, the inclusion of interstitial- substitutional interactive terms, such as N Mn and Ti C, improved FN predictive accuracy. From

a metallurgical viewpoint, elements such as titanium and carbon can interact

to form a new phase by the following reaction:

Ti + C = Ti C. (22)

The equilibrium constant for this reaction can be represented by:

K = f Ti Cl (23)

[Ti] [G]

where the brackets represent the concentrations of the various components.

This indicates that the product of the concentrations of these two elements

is directly related to the formation of this phase, and if this phase occurs

in stainless steel welds, such an interaction term should be included in an

FN predictive equation. Similar interactive terms would be expected for

other reactions, intermetallic compounds, and ordered structures.

In ferritic phases, these types of interactions have been confirmed in

metallurgical studies. 59 Internal friction studies in commercial steels

have shown that carbon and nitrogen interstitial atoms hold manganese,

titanium, and other substitutional atoms in close association. 60 Clusters

of carbon and nitrogen have been observed about titanium atoms in Fe—0.15 Ti

alloys by a field ion microscope atom probe. 61 The improvements in FN prediction found in this study by the inclusion of interstitial - subs titutional

interactive terms implies that similar associations are present in stainless

steel alloys. Suspected interactions between nitrogen and manganese in the

mechanical portion of this study support this premise.

In the case of the titanium—carbon interactive term, the precipitation

of titanium carbides must be considered in the FN predictive equations. If some portion of the titanium and carbon is removed from their respective

roles during solidification by formation of a precipitate, then the effects

of these elements on the ferrite content must also be reduced.

Interactions between unlike substitutional atoms were also indicated in

this study. Most notably, the chromium—molybdenum interaction. It is difficult to speculate on the specific significance of this interaction. The

nonlinear trends observed in figure 46, however, indicate that the interaction has a pronounced effect on FN.

Although the interactive terms developed in this study are not proposed

for use at this time, they may identify some of the more important interactions occurring between the alloying elements in stainless steels. Until

all the pertinent linear and nonlinear contributions made by these alloying

elements are determined and incorporated into a fundamental equation form

such as equation 5, more specific information regarding the physical significance of these equations will be unobtainable.

213
V. CONCLUSIONS

1. The yield strength at 4 K for welds evaluated in this study can be

expressed by the equation:

Yield Strength (MPa) = 334 + 3407N + 15Mn + 46Mo (7)

2. The notch ductility at 76 K for the welds evaluated in this study was found to be principally a function of strength:

LE (nun) = 0.95 - 6 x 10' 4 (Yield Strength, MPa) (10)

3. Molybdenum additions approaching 4 wt.% can severely reduce toughness in

type-18-8 nitrogen—manganese stainless steel weld metals.

4. The DeLong diagram was quite accurate in predicting FN for type- 308 and

316 stainless steel compositions, but had to have an average error near

2 FN for type- 309 alloys.

5. An improved predictive equation was developed having chromium and nickel

equivalents similar in complexity and form to the DeLong equivalents.

FN1 = -15.2 + 2 . 2

[ Cr + Mo + 0.5Si + 15Ti + 0.8Cb] - 1.9[Ni + 17N + 30C]

.

(14)

This equation predicted FN better than the DeLong diagram for common

alloy compositions. The equation was substantially better than the

DeLong diagram in predicting the FN of alloys with extended manganese and nitrogen contents.

6. The accuracy of FN prediction can be improved further by investigating

the following topics:

a. The grouping of welds by primary solidification modes prior

to FN equation development

b. The development of alternative, nonlinear equation forms

c. The identification of important interaction terms to be

included in predictive equations

d. The adjustment of principal compositional coefficients,

particularly for titanium

214
VI . ACKNOWLEDGMENTS

I am indebted to many people for guidance, support, and tolerance during

the research for and preparation of this manuscript.

For financial support, I acknowledge the United States Department of

Commerce, the United States Department of Energy, and the Welding Research

Council Subcommittee on Stainless Steel Welding.

For keeping the study on track, I owe much to the guidance of Tom

Siewert, Dave Olson, and Dave Matlock.

For tolerance, above and beyond the call of duty, I owe special thanks

to: my wife Linda, my daughter Kimberlie, my parents, coworker Bob Walsh,

and typist Chris King.

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Corrosion Testing (E 23-72), ASTM, 273-293 (1980).

56. McCowan, C. N. , Siewert, T. A., Kivineva, E. , "Molybdenum in Stainless

Steel SMA Welds for Cryogenic Service," Presentation AIME Welding

Metallurgy of Structural Steels, Denver, Colorado, February 23—26,

(1987) , in press

57. Purtscher, P. T. , Reed, R. P., Advances in Cryogenic Engineering

Materials, Plenum Press, New York, To be published (1987).

58. Reed, R. P., "Martensitic Phase Transformations," Chapter 9 in

Materials at Low Temperatures, American Society for Metals, Metals

Park, Ohio, 303-307 (1983).

59. Leslie, W. C., The Physical Metallurgy of Steels, McGraw-Hill,

New York, 136-139 (1981).

60. Dijkstra, L. J., Sladek, R. J., Transactions of the AIME 69, 197

(1953)

61. Hasson, D. T. , Arsenault, R. J., Treatise on Materials Science and

Technology, Academic Press, New York, 179 (1972).

219
Table 1 . Various Coefficients for Manganese and Nitrogen

Determined in Previous Studies

Researcher

Coefficient

or Constant

N Mn

Principal

Matrix Variations

(wt.%)

Comments

Schaeffler [33] 0.5 Mn Ni (0 to 30) 308, 309, 310, 316, 410

and 502 stainless steels

DeLong [35]

Hull [43]

Espy [37]

Hammar

and

Svenson [44]

Mel'Kumor

and

Topilin [45]

Szumachowski

and

Kotecki [36]

Okagawa

et al. [46]

30 N N (0.03 to 0.22) 309, 308L, 316, 316L, 347 (low Mn) stainless steels

18.4 N 0.11 Mn - Ni (0 to 22) 70 chill-cast stainless

0.0086 Mn2 Mn (0 to 20) steel alloys. Mo, Si, V, W, N (0 to 0.15) Ti, Cb, Ta, Ai , Co, Cu, C also varied. Cr - 14 to 20 wt.%

30 (N-0.045) 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

20 (N-0.45) systematically. The 3 N coefficients are for < 0.20,

- < 0.25, and < 0.35 wt.% N, respectively

.

14.2 N 0.31 Ni (9.0 to 14) 130 austenitic alloys

N (0.01 to 0.20) (ingots). Mn content - 1 to

C (0.04 to 0.10) 2 wt.% Cr - 17 to 25 wt.%

20 N - Ni (10 toi 16) Austenitic stainless steel.

N (0.05 to 0.46) 0.5 Mn, 25 Cr, 0.09 C

- 0.35 Mn (1 to 12) 308, 309L, 316L, 307 stainless steels. Low nitrogen (0.05 wt.%)

13.4 N -

N (0.04 to 0.29) GTA weld (no filler metal) on 304L stainless steel base

metal. N added through

shielding gas.

220
Table 2. The SMA Weld Metal Compositions for the Manganese—Nitrogen Test Matrix.

Alloy

No. c Mn Si P

1 s Cr Ni Mo N B Pad

FN

Plate+ 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 .034 1 .49 0 .34 0 .020 0 .006 15 .37 9 .06 2 .03 0 .034 # 3 .4 4 .4 2 .2

3 0 .034 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 .54 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 .013 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..017 0,.009 17,.71 9..10 0..02 0,.106 # 1..5 2..7 -12..16

11 0. 031 3..18 0 .34 0..015 0..006 18..99 9..95 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 14 0. 038 9.,38 0..35 0.,017 0.,009 18.,98 9.,14 0..02 0.,163 # 1..7 2 . 8 -12 .,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.,252 # 1..2 1. 8 2 . 83

17 0. 031 2. 90 0..30 0. 013 0. 006 18. 79 9. 12 0.,02 0.,104 2. 0 3. 8 2 . 83

18 0. 033 3. 25 Or; 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. 013 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.

Repeated welds.

+ FN plate is the average of FN bead surface and FN impact specimen surface.

*Not determined.

221
Table 3. The Range of Composition for Weld Metals

of the FN Test Matrix

Element

Range of Composition, wt . %

Min

.

Max

.

Mean

Standard

Deviation

C 0.012 0.200 0.044 0.019

Mn 0.350 12.67 2.317 2.124

Si 0.150 0.940 0.456 0.147

Cr 14.74 26.90 20.31 2.360

Ni 4.610 33.50 11.27 2.286

Mo 0.010 6.850 1.102 1.381

N 0.019 0.300 0.080 0.045

Cb k 0.880 0.030 0.106

Ti k 0.540 0.020 0.034

*Not determined

Table 4. -Definitions of the Notations Used to Identify the

Previously Developed Ferrite -Number Predictive Equations

Equation

Abbreviation Description

FNCD FN calculated using the DeLong equation form

FNCD = - 30.7 + 3 . 5 [Cr + Mo + 1.5(Si)] - 2 . 5 [Ni + 30 (n) + 30(C) + 0.5(Mn)]

FNCDK FN calculated using the DeLong equation form

but with the Szumachowski—Kotecki 0.34 constant

substituted for the 0.5(Mn) term in the nickel

equivalent

FNCH FN calculated using the Hull equation form,

FNCH = +2 . 88 [ Cr + 1.1 (Mo) + 0.5(Si)] - 3 . 13 [Ni + 18.6 (N) + 24.5(C) + O.ll(Mn) - 0 . 0086 (Mn2

)

]

222
Table 5. Mechanical Property Test Results

Alloy 4-K Tensile Properties 76 -K CVN Properties

°y i MPa

<?T > MPa

Elongation,

%

Reduction,

% LE,f

mm

AE,t

J

1 469 1586 26* 16.6 0.686 46

2 527 1570 20 13.5 0.711 39

3 504 1432 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

11 932 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

14 1075 1527 28.5 26.6 0.305 34

15 1344 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 1125 11.3 9.6 0.432 41

22 1250 1498 15.5 15.4 0.152 21

^Fractured outside the gauge marks

|LE = lateral expansion

fAE = absorbed energy

Table 6. The Ductile Dimple Densities of the Fracture Surfaces

Alloy Specimen Type

Ductile Dimple Density

per mm2

1 4-K tensile 320,000

1 76 -K Charpy 410,000

6 4-K tensile 440,000

6 76 -K Charpy 440,000

21 4-K tensile 350,000

21 76 -K Charpy 460,000

223
Table 7. Ferrite -Number Predictive Equations

Equation

Number Equation Statistics

Cases F R

2 STD

GROUP 2

11 FN , 4.48 + 0. 34[FNCD] 746 582 0.44 2.5

12 FN = 2.43 + 0 . 51 [ FNCDK] 746 930 0.55 2.2

13 FN = -5.91 + 0 66 [ FNCH] 746 869 0.54 2.2

14 FN1 = -15.2 + 2 2 [ Cr + Mo + 15 (Ti) + 0 . 8 ( Cb

)

+ 0 . 5 (Si)

]

- 1 . 9 [Ni +17 (N) +30(C)

]

746 1454 0.66 1.9

GROUP 3

15 FN = 1.8 + 0 . 65 [FNCD] 582 999 0.63 1.8

16 FN = 0.12 + 0 . 98 [ FN1

]

582 1143 0.66 1.7

GROUP 4

17 ~FN2 _ -14.6 + 2 4 [ Cr + Mo + 15 (Ti) 1215 0.67 1.7 + 0 . 8 (Cb) + 0 . 7 (Si)

]

- 2 . 2 [Ni + 23 (N) +27(C)]

NONLINEAR

18 FN3 2 . 79 [ Cr + 0 . 7 (Mo) + 15 (Ti) 746 2021 0.73 1.7 + 0 . 8 ( Cb ) ] - 473 (Ti • C) - 0 . 002 (Cr -Mo) 2 + 0 . 0002 (Cr) 2 (Mo) 3 - 2 . 36 [Ni + 31(C) - 22 (N) - 0 . 05 (Mn)

]

+ 1 . 6 (Ni) .C + 4(N) + 0 . 1 (Mn)

]

- 7 . 9 (Mn) [N - 0 . 2(Mn-N)

]

2 + 0 . 5 (Mn)'

19 FN4 /, q _i_ TO 2 . 3 ( Cr") + 2.4(Mo') + Si + 2 (Cb) + 33 (Ti) + 32'

40 -+* jZ

2 (Ni) + 57(C) + 29 (N) + 47

224
Table 8. The Szumachowski—Reid Weld Compositions 2

Alloy C Mn Si Cr Ni Mo N

AE LE

T7NT J ft-lbf /xm mil

r JN

a 0.05 2.21 0.35 15.2 9.1 1.66 0.04 36.6 27.0 610 24.0 2.0

b 0.05 2.17 0.25 15.3 9.2 1.67 0.04 35.9 26.5 526 20.7 2.0

c 0.06 1.97 0.36 16.7 8.7 1.67 0.04 32.5 24.0 457 18.0 5.8

d 0.05 2.23 0.41 15.3 9.2 1.70 0.08 28.6 21.1 312 12.3 0.8

e 0.05 2.00 0.43 17.8 8.7 1.72 0.08 30.2 22.3 271 10.7 3.2

f 0.04 2.15 0.35 15.3 9.4 1.73 0.08 31.6 23.3 305 12.0 0.6

g 0.06 1.62 0.38 19.5 9.8 0.20 0.06 18.4 13.6 160 6.3 4.0

h 0.06 2.26 0.36 21.1 9.5 0.20 0.08 12.7 9.4 84 3.3 0.0

Table 9. The Unfactored Coefficients for the FN Predictive Equations

DeLong 3.5 3.5 5.2 1.7 — -2.5 -75 -75 -1.3

FNCH 1.9 2.1 0.9 — — -2.1 -38 -50 +0 . 2(Mn) - 0 . 018(Mn)"

FN1 2.2 2.2 1.1 1.8 33 -1.9 -32 -65 —

FN2 2.4 2.4 1.6 1.9 35 -2.2 -50 -60 —

225
NICKEL

EQUIVALENT

=

y.Ni +

30X

y.C +

30X

y.N +

0.3 X

y.Mn

Figure 1. The Schaeffler constitutional diagram for stainless steel weld metal. 33

226
O 2 4 6 8 IO 12 14 16 18 20 22 24 26 28 JO 32 J4 J6 38 40 Chromium Equivalent' -% Cr + % Mo t 1.5 x % Si + 05 x % Cb

Figure 2. The DeLong constitutional diagram for stainless steel weld metal. 34 - 35

227
A

Gr

Ni, wt.%

Figure 3. Isothermal diagrams at 1000°C from the Fe—Cr—Ni ternary diagram.

228
Figure 4. Stress-strain behavior of an 18Cr— 8Ni austenitic

stainless steel as a function of temperature. 60

1.6

m

Ph° 1.2

to

COw

o£H

LO 0.8

0.4

0 0.1 0.2 0.3 0.4 0.5

STRAIN

Figure 5. Stress-strain behavior of an 18Cr-8Ni austenitic

stainless steel at 4 K as a function of alloy content. 60

229
~3

Figure 6

E

Charpy V-notch absorbed energy versus Lateral expansion

measurements in materials of varying strength levels. 52

230
Figure 7. The 4-K tensile testing apparatus.

231

mur
Stainless Steel

Figure 8

.

Detail of the load train and specimen configuration

during tensile testing.

232
Figure 9. Detail of the extensometer configuration

0.01 mm

Figure 10. Optical micrograph showing inclusions on a mechanically

polished cross section of weld 1. This inclusion morphology

and distribution is representative of all the welds.

233
Figure 11. Optical micrograph of the microstructure of weld 6.

Vermicular and lacy ferrite morphologies were commonly

observed. A hydrochloric-nitric-acetic acid etch was used.

%

l I 0.04mm

Figure 12. Optical micrograph of the microstructure of weld 6.

The weld was etched in a potassium hydroxide solution.

234
1 I 0.2 mm

i i 0.04 mm

Figure 13. Optical micrograph of the microstructure near

an interpass region in weld 22. The weld was etched in a 33 percent nital solution.

233
Number

of

inclusions

per

mm

Figure 14. Inclusion distributions found for four of the weld deposits.

Weld 6 had more small diameter inclusions than the other welds.

Figure 15. SEM micrograph showing the inclusion

morphology observed for alloy 15.

236
Figure 16. SEM micrograph showing the inclusion

morphology observed for alloy 6.

Figure 17. Stress-strain curves characterizing the

deformation behavior of the welds.

237
True Strain, %

Figure 18. Serrated stress -strain curve characteristic

of the welds tested at 4 K.

1400

1300

1200

. 1100

- 1000

"So 900

S 800

ro Cl

CO 700

600

500

400

• 4 K Yield Strength

0 4 K Tensile Strength

240

220

200

180

160

140 *

120

100

80

60

40

0 08 0.12 0.16 0.20 Nitrogen, wt. %

0.24 0.28

Figure 19. Strength versus nitrogen content of the welds

238
200

180

160

140 --

120

"*

100

80

60

0 1 2 3 4 5 6 7 8 9 10

Manganese, wt. %

1400 - i i I i i—rn i

i r 1300

1200 V

1100 V

V

1000

V

900

V

V V • •

800 — •

700 •

o o

600 — o

500 _ o

o

<8> o o 0.05 Wt. % N

400 • 0.10 wt. % N

V 0.15 wt. % N

| | 1 1 1 1 1

1 1 0.25 wt. % N

Figure 20. Yield strength versus manganese content of the welds.

-> 65

jjjj 60

I 55 - 50

45

§ 40

5 35

g 30

°

25

S 20 ^ 15

10

5

0 0.04 0.08 0.12 0.16 0.20 0.24 0.28

Nitrogen, wt. %

Figure 21. Charpy V-notch absorbed energy versus nitrogen

content in the welds.

— T 1 1 1 n 50

Alloy Mn wwt.%

1.7 1.5 — 45 o 11,17,18 3 a 4.9,12 6 40 — 35

DO A o

AO

eg 30

*0^ O A <6

-

25

20

o — o° 15

o — 10

5

1 J L 1 1

I 1

239
070 0028

0024

0.020

0016

0.012

0 008

0004

c

Nitrogen, wt. %

Figure 22.- Lateral expansion versus nitrogen content in the welds.

in

0.010 0.020 0.030 in

LOXL

Figure 23. The relationship between 4-K yield

strength and 76-K lateral expansion.

240
Alloy 1

low nitrogen

Alloy 22

high nitrogen

Figure 24. The effect of nitrogen on the fracture surface

appearance of Charpy V-notch specimens.

241
Figure 25. Similarities in the shape and spacing of ferrite

morphologies of polished specimens and the tearridge networks on the fracture surface for weld 12.

242


Figure 26. Similarities in the shape and spacing of ferrite

morphologies of polished specimens and the tearridge networks on thefracture surface for weld 12.

243
Figure 27. Fractograph of Charpy V-notch specimen, taken near the notch

of weld 6. The ductile dimple density shown here is similar

to those observed on tensile specimen fracture surfaces.

10 urn

Figure 28. Fractograph showing one of the large inclusions found

in the 3.84 wt. pet. molybdenum alloy (weld 6).

244
Figure 29. Fractograph taken near the center of a weld 1 tensile

specimen. It is representative of the dimple size and

morphologies commonly observed.

l_

500 ym

j

Figure 30. Fracture surface features on a Charpy V-notch specimen

near the notch of a weld 1 specimen. Similar markings

were observed in other welds.

245
Figure 31. Fracture surface features observed near the center

of a weld 13 Charpy V-notch specimen. Similar features

were observed in this area for the other welds.

246
I I 10 pm

Figure 32. Isolated ferritic area of a weld metal tensile

specimen that fractured by cleavage.

| ) 20 pm

Figure 33. Isolated ferritic area of a weld metal tensile

specimen that fractured by cleavage.

247
>>

4-> •Hm

C &*

<u jjC cn

a>>

•H

4-> ca

<u OS

m

C

QJ 4JC cn

aj no > rH

•H 0) U CS

CO

l"“) <u OS

20, deg.

Figure 34. X-ray determination of the phases present in welds 6 and 13.

The principal phase present in the welds was austenite.

248
ft- lb

5 10 15 20 25 30 35 40 45 50

76K CVN Absorbed Energy, J

Figure 35. The relationship between lateral expansion and

absorbed energy for the Szumakowski-Reid data. 27

249
ft- lb

5 10 15 20 25 30 35 40 45 50

Figure 36. The relationship between lateral expansion and

absorbed energy for the welds in this study.

250
Cr eq

Figure 37. Predicted error of FN for data group 1 using

the DeLong-type equation (FNCD)
Figure 38. Predicted error of FN for data group 1 using the 0.34 constant

for the manganese term in the DeLong type equation (FNCDK)

.

252
Cr eq

Figure 39. Predicted error of FN for data group 1 using the FN1 equation.

253
24n

20'

FN Calculated - FN Measured

Figure 40. The frequency (%) versus FN predicted error for group 1 when the FNCD, FNCDK, and FN1 equations were used.

254
FN

Calculated

(FNCDK)

— FN

Measured

20

16 17 18 19 20 21 22 23 24 25 26 27 28 29 30

wt. % Cr

Figure 41 . FN predicted error versus chromium content for data group 1 using the 0.34 constant for the manganese term in the DeLongtype equation (FNCDK)

255
FN

Calculated

(

FN

I)

—

FN

Measured

16 17 18 19 20 21 22 23 24 25 26 27 28 29 30

wt. % Cr

Figure 42. FN predicted error versus chromium content for

data group 1 when the FN1 equation was used.

256
wt. °/o Mo

Figure 43. FN predicted error versus molybdenum content for

data group 1 when the FNCD equation was used.

257
20

*o

CD

i—

13CO

CO

CD

15

10

5

2

u_

0

OCD -5

Jo

D

O -10

CO O

Z

LL

-15

-20

fccPo a

jo o

0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

wt. % Mo

Figure 44. FN predicted error versus molybdenum content for

data group 1 when the FN1 equation was used.

258
Figure 45. FN predicted error versus molybdenum content for

data group 2 when the FN1 equation was used.

259
FN

Calculated

(

FNI )

—

FN

Measured

20

15

10

5

u

D, ftp u 0 mm

-5

IF

-10

-15 —

-20

u

mi

."rfffi

Q ^ ", Qc

q “g3 y a

U

rf

„

°D

Q O ^S> a a u

cP a a u

0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0

wt. °/o Mo

Figure 46. FN predicted error versus molybdenum content for

weld data having FN values between 1 and 30 (FNI).

260
20

o

0)

3

in

05

CD

LL

Q

o

O

0)

13 _o

03 O

2

LL

15

10

5

0

-5

-10

-15

-20

?! a s° Q D

o

__ n qO a q

u du

a

0.05 0.10 0.15 0.20

wt. % N

0.25 0.30 0.35

Figure 47. FN predicted error versus nitrogen content for

data group 1 when the FNCDK equation was used.

261
20

-15

-20

0.05 0.10 0.15 0.20

wt. % N

0.25 0.30 0.35

Figure 48. FN predicted error versus nitrogen content for

data group 1 when the FN1 equation was used.

262
20

oCD

13CO

03

CD

~o

0)

_o

cn O

-15

-20

0.05 0.10 0.15 0.20 0.25 0.30 0.35

wt. % N

Figure 49. FN predicted error versus nitrogen content for

welds having FN values between 1 and 30 (FN1).

263
20

wt. % N

Figure 50. FN predicted error versus nitrogen

data group 2 when the FN1 equation

0.30 0.35

content for was used.

264
20

o

-20 1 1 L1 2 3 4 5 6 7 8 9 10 11 12 13 14

wt. % Mn

Figure 51. FN predicted error versus manganese content for

data group 1 when the FNCD equation was used.

265
20

"O0

3

C/) 0

o

o

-*— 0

0

3

O

0O

8 9 10 11 12 13 14

wt. % Mn

Figure 52. Predicted error versus manganese content for

data group 1 when the FN1 equation was used.

266
FN Calculated - FN Measured

Figure 53. The frequency (%) versus FN predicted error

for type 308 stainless steels.

FN Calculated - FN Measured

Figure 54. The frequency (%) versus FN predicted error

for type 309 stainless steels.

267
30 -I 28 -

26 - 24

1

22- 20

18-

a>o- I0-1 a> £ 8

-

6

- 4 ^

2

-

316

FNI

FNCD

-10 -8 ' -6 -4 -2 6 2 4 6 8 10 FN Calculated - FN Measured

Figure 55. The frequency (%) versus FN predicted error

for type 316 stainless steels.

Figure 56. The manganese-nitrogen interaction.

268
Appendix I. FN Data Base

The weld metal compositions and FN values collected for the data base

to develop and evaluate equations predicting ferrite content.

c Mn Si Cr Ni N Mo Cb Ti FN

.04 1 .54 .50 18.00 10.80 .16 .13 .00 .00 .2 .02 1 .49 .46 18.20 10.40 .49 .21 .00 .00 .2 .05 1 .43 .45 16.70 10.20 .15 .05 .00 .00 .3 .04 1.55 .60 18.40 11.70 .15 .04 .00 .00 2.2

.02 1 .51 .47 18.60 9.10 .09 .22 .00 .03 .5 .04 1 .54 .50 18.00 10.80 .16 .04 .00 .00 2.4

.02 1 .49 .46 18.20 10.40 .49 .10 .00 .00 .5 .02 1 .53 .41 18.30 9.10 .31 .18 .00 .00 .4 .02 1 .65 .48 18.40 9.20 .36 .17 .00 .01 3.2

.02 1 .53 .41 18.30 9.10 .31 .12 .00 .00 1 .6 .03 1 .61 .35 18.00 9.40 .13 .06 .00 .00 2.6

.02 1 .51 .47 18.60 9.10 .09 .12 .00 .03 1 .9 .03 1.72 .57 17.50 13.40 2.72 .25 .00 .01 .2 .03 1 .72 .57 17.50 13.40 2.72 . 14 .00 .01 .2 .04 1 .06 .78 18.60 14.50 2.07 .02 .02 .02 .5 .03 1 .53 .47 17.40 11.20 2.71 .19 .00 .00 .3 .03 1 .53 .47 17.40 11.20 2.71 .14 .00 .00 2.0

.05 1 .73 .36 17.20 11.80 2.78 .04 .00 .00 3.7

.02 1 .63 .54 17.30 10.40 2.81 .20 .00 .00 .6 .04 1 .64 .40 17.10 10.80 2.50 .06 .00 .00 5.2

.02 1.63 .54 17.30 10.40 2.81 .12 .00 .00 3.5

.03 1 .90 .61 18.40 14.10 3.20 .05 .00 .00 1 .8 .03 1 .34 .43 18.40 14.20 3.21 .03 .00 .00 2.5

.01 1 .55 .46 19.50 24.80 4.45 .06 .00 .01 .2 .02 1 .40 .50 19.60 24.80 4.33 .04 .00 .00 .2 .02 1 .59 .51 16.80 14.50 4.43 .30 .00 .02 .3 .02 1 .59 .51 16.80 14.50 4.43 . 1 1 .00 .02 .2 .03 1 .47 .40 17.00 15.50 4.47 .03 .00 .01 .6 .02 1 .41 .50 17.60 13.50 4.58 .20 .00 .00 .2 .02 1 .49 .49 17.00 14.70 4.57 .06 .00 .02 1 . 1 .02 1 .41 .50 17.60 13.50 4.58 .10 .00 .00 1 .9 .08 5.79 .64 26.10 20.00 .57 .05 .02 .02 .5 .04 6.33 .37 18.00 15.00 2.61 .05 .01 .04 .2 .07 7.77 .43 17.20 9.30 .34 .04 .00 .04 2.4

.01 6.20 .37 17.90 7.90 .05 .08 .00 .00 1 . 1 .05 1 .63 .46 17.10 10.70 2.13 .23 .00 .00 . 1 .05 1 .47 .48 21 .30 11.50 .17 .24 .00 .00 .3 .05 1 .50 .48 17.80 9.40 .17 .18 .00 .00 . 1 .04 1 .59 .54 23.80 13.50 .10 .20 .00 .00 .5 .04 1 .48 .55 26.90 15.20 . 13 .24 .00 .00 3.7

.04 1 .33 .43 16.80 12.90 2.71 .06 .00 .00 .3 .04 1 .42 .46 17.00 13.00 2.71 .06 .00 .00 .9 .05 1 .63 .46 17.10 10.70 2.13 .09 .00 .00 1 . 6 .04 1 .59 .54 23.80 13.50 . 10 .05 .00 . 00 11.6

.05 1 .50 .48 17.80 9.40 .17 .05 .00 .00 3.0

.04 1 .48 .55 26.90 15.20 .13 .05 .00 . 00 19.3

.05 1 .47 .48 21.30 11.50 .17 .05 .00 .00 8.4

.03 1 .42 .46 16.70 10.70 2.70 .05 .00 .00 5.5

.06 1 .62 .38 19.53 9.84 .20 .06 .00 .00 4 . 0

.05 1 .65 .35 18.81 10.40 .20 .06 .00 .00 1 .8 .05 1.19 .35 19.69 10.02 .13 .04 .00 .00 7.7

.05 1 .19 .32 19.13 10.35 .12 .04 .00 .00 5.4

.05 1 .13 .41 18.91 11.29 .13 .04 .00 .00 . 1 .04 2.27 .31 18.97 12.30 2.28 .03 .00 .00 7.2

.04 2.32 .36 18.37 12.50 2.29 .03 .00 .00 4.8

.04 2.37 .33 18.08 13.52 2.22 .03 .00 .00 .8 .04 2.27 .38 17.53 14.25 2.23 .03 .00 .00 . 1 .04 2.26 .37 17.36 16.70 2.15 .03 .00 .00 . 1 .03 2.54 .27 17.82 11.98 2.28 .06 .00 .00 5.4

.03 2.52 .35 17.81 12.61 2.19 .06 .00 . 00 1 .8

269
Appendix I . FN Data Base , continued

C Mn Si Cr Ni N Mo Cb Ti FN

.04 2 .51 .30 17 .80 12 .90 2 .25 .05 .00 .00 .4 .03 2 .64 .32 17 .74 15 .60 2 .12 .05 .00 .00 .0 .03 2 ,44 .30 18 .22 1

1

.63 2 .29 .09 .00 .00 1 .8 .03 2 .42 .34 17 .91 12 .50 2 .18 .10 .00 .00 .2 .03 2 .42 .29 17 .57 13 .03 2 .18 .10 .00 .00 .0 .03 2 .46 .31 17 .36 15 .48 2 .14 .09 .00 .00 .0 .05 1 .16 .40 17 .70 8 .80 1 .69 .04 .00 .00 9 .9 .06 1 .97 .36 16 .68 8 .74 1 .69 .04 .00 .00 5 .8 .05 2 .21 .35 15 .24 9 .09 1 .66 .04 .00 .00 1 .5 .05 2 .17 .25 15 .31 9 .19 1 .67 .04 .00 .00 1 .4 .05 2 .18 .30 15 .00 1

1

.47 1 .60 .03 .00 .00 0 .05 1 .30 .43 17 .82 8 .75 1 .69 .07 .00 .00 8 .2 .05 2 .00 .43 17 .82 8 .72 1 .72 .08 .00 .00 3 . 1 .05 2 .23 .41 15 .33 9 .19 1 .70 .08 .00 .00 1 .4 .04 2 .15 .35 15 .28 9 .36 1 .73 .08 .00 .00 .6 .05 2 .10 .31 14 .93 1

1

.52 1 .64 .07 .00 .00 .3 .03 1 .12 .33 21 .60 1

1

.45 .04 .05 .00 .00 1

1

. 1 .03 2 .16 .36 21 .81 1

1

. 60 .05 .06 .00 .00 10 . 4 .03 3 .24 .34 21 .78 1

1

.67 .05 .06 .00 .00 10 .6 .04 4 .17 .36 21 .87 1

1

.49 .01 .06 .00 .00 1

1

.5 .04 5 .36 .37 21 .99 1

1

.56 .01 .06 .00 .00 1

1

.3 .04 5 .81 .35 21 .99 1

1

.76 .01 .06 .00 .00 1

1

.3 .04 6 .20 .35 22 .20 1

1

.65 .06 .06 .00 .00 1

1

.5 .04 8 .09 .36 22 .20 1

1

.75 .06 .06 .00 .00 1

1

4 .04 8 .57 .32 22 .19 1

1

.85 .05 .06 .00 .00 10 9 .04 10 .04 .34 22 .28 1

1

.88 .01 .06 .00 .00 12 6 .04 1

1

.09 .36 22 .26 12 .00 .01 .07 .00 .00 12 .0 .04 12 .67 .33 22 .39 12 .18 .02 .07 .00 .00 12 . 1 .07 1 .50 .27 23 .08 12 .28 .09 .07 .00 . 00 8 8 .06 2 .00 .27 23 .19 12 .18 .09 .07 .00 .00 8 .3 .07 3 .59 .24 23 .32 12 .41 .09 .07 .00 .00 9 .2 .07 4 .76 .27 23 .42 12 .41 .09 .07 .00 .00 9 .6 .06 ' 5 .61 .28 21 .92 1

1

.51 .06 .07 .00 .00 7 .6 .07 6 .17 .21 21 .95 1

1

.60 .08 .08 .00 .00 7 .3 .07 9 .28 .23 22 . 1

1

1

1

.74 .08 .07 .00 . 00 8 0

.08 10 .85 .23 22 . 17 1

1

.82 .08 .08 .00 .00 8 .2 .07 4 .92 .23 21 . 77 1

1

. 40 .13 .07 .00 .00 7 .6 .06 3 , 92 .24 21 .83 1

1

.63 ,13 .07 .00 .00 7 .0 .07 3 ..19 .25 21 .64 1

1

.52 . 13 .07 .00 .00 6 .5 .06 2 ..20 .25 21 ..55 1

1

..30 13 .06 .00 .00 6 .5 .06 1 . 40 .27 21 . 61 1

1

. 60 14 .06 .00 .00 6 ,.0 .04 1 . 15 .41 17 . 85 1

1

. 80 2 . 18 .03 .00 .00 7 ,. 1 .04 2 . 07 .39 17 . 98 1

1

. 55 2 . 12 .03 .00 . 00 6 . 9 .03 3 . 03 .39 18 . 1

1

1

1

. 45 2 . 20 . 04 .00 .00 6 . 9 .05 3 . 94 .41 18 . 1

1

1

1

. 80 2 . 16 .03 .00 .00 6 . 8 .05 5 . 13 . 41 18 . 20 1

1

. 68 2 . 16 . 04 .00 .00 7 . 1 .04 7 . 44 .39 18 . 19 1 1 . 75 2 . 24 .04 .00 .00 6 . 6 .04 8 . 52 .42 18 . 30 1

1

. 73 2 . 26 .04 .00 .00 7 . 8 .05 9 . 48 .40 18 . 33 1

1

. 80 2 . 18 .05 .00 .00 6 . 9 .05 10 . 74 .40 18 . 28 1

1

. 88 2 . 20 .04 .00 .00 5 . 4 .05 12 . 39 .39 18 . 40 1

1

. 90 2 . 20 .05 .00 .00 2 . 6 .05 1 . 33 .22 19 . 90 9 . 49 07 .03 .00 .00 1

1

. 2 .05 1 . 77 .24 20 . 38 9 . 65 07 .03 .00 .00 12 . 2 .05 2 . 36 .25 20 . 98 9 . 52 07 . 1

1

.00 .00 7 . 9 .09 3 . 60 .47 19 . 97 9 . 49 1 . 10 .06 .00 .00 9 . 9 .08 3 . 40 .35 19 . 95 9 . 51 1 . 18 .07 .00 .00 10 . 1 .08 3 . 41 .37 19 . 77 9 . 18 1 . 1

1

.07 .00 .00 10 . 1

. 04 1 . 92 .52 19 . 20 13 . 20 2 . 20 .06 .00 .04 4 . 0 .13 2 . 15 .54 26 . 40 20 . 50 22 .06 .00 .05 0 .06 1 . 95 .59 19 . 80 9 . 50 09 .04 .00 .01 8 . 0

.03 1 . 81 .50 19 . 20 9 . 50 15 .07 .00 .01 6 . 0 .03 1 . 68 .42 19 . 40 9 . 50 13 .09 .00 .03 4 . 0 .04 1 . 73 .49 20 . 00 9 . 90 12 .07 .00 .03 6 . 0

270
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.04 1 .99 .64 19 .90 9 .70 .43 .06 .00 .01 6.0

.04 2.14 .31 19 .60 9 .70 .44 .06 .00 .01 6.0

.05 2.63 .23 19 .90 9 .80 .44 .05 .00 .01 6.0

.05 1 .78 .58 19 .80 9 .60 . 1

.07 .00 .01 6.0

.03 1 .52 .49 18 .90 9 .70 .15 .09 .00 .01 4.0

.20 1 .83 .49 25 .90 20 .40 .21 .07 .00 .02 .0

. 11 .62 .86 25 . 10 19 .70 2 .60 .05 .00 .01 5.0

.05 1 .56 .51 18 .00 13 .30 2 .20 .08 .00 .00 .0 .03 1 .57 .54 17 .90 13 .50 2 .20 .05 .00 .00 .0 .20 1 .63 .45 15 . 10 33 .50 .72 .04 .00 .03 .0 .06 1 .93 .47 19 .90 9 .70 .13 .06 .00 .02 6.0

.03 .90 .71 20 .40 10 .00 .02 .09 .00 .03 12.0

.03 2.20 .26 17 .92 13 .35 2 .16 .03 .04 .06 .9 .04 2.15 .28 18 .00 13 .55 2 .12 .02 .04 .06 .7 .06 1.11 .36 20 .06 9 .90 .14 .05 .02 .05 6.6

.03 2.22 .38 17 .79 13 .18 2 .14 .03 .04 .06 1 .2 .03 2.12 .37 18 .13 13 .48 2 .19 .03 .04 .06 1 .0 .03 2.13 .29 18 .01 13 .15 2 .16 .02 .04 .07 1 .7 .03 2.12 .31 18 .12 13 .42 2 . 18 .02 .04 .08 1 .5 .04 5.82 .71 21 .08 10 .15 1 .74 .21 .03 .04 11.9

.04 1 .80 .46 18 .09 13 .60 3 .58 .08 .03 .03 4.4

.04 11.20 .25 17 .60 5 .21 .12 .17 .02 .02 4.8

.02 1 .08 .33 19 .69 9 .81 .12 .06 .02 .04 7.3

.08 1 .00 .39 24 . 1

12 .86 .19 .07 .03 .05 9.2

.02 1 .02 .78 19 .56 1

,.63 2 .20 .05 .04 .06 9.8

.03 .88 .59 18 .62 1

.53 2 .00 .07 .03 .05 6.5

.02 .94 .62 20 .00 10 .12 .12 .03 .03 .05 11.1

.05 1 .00 .68 20 .06 10 .00 .14 .05 .03 .06 7.8

.05 .98 .68 20 .37 9 .96 .13 .06 .03 .05 8.8

.03 1 .13 .39 20 .47 9 ,.97 .08 .08 .03 . 04 8.8

.06 1 .05 .39 19 ..84 9 ,.60 . 1

.06 .03 .05 7.8

.03 .94 .43 20 ,.85 9 .66 .16 .05 .03 .05 11.9

.03 1 .04 .48 20 .69 9 .60 .20 .08 .03 .05 11.4

.03 .96 .38 19 ,.88 9 .52 .18 .07 .03 .05 8.5

.04 1 .74 .40 18 ..31 12 .47 2 .24 .05 .04 .05 6.4

.03 .93 .33 18 ..78 11 .,71 2 .25 .07 .03 .05 8.3

.04 .91 .33 18 ..89 1

..70 2 .30 .06 .03 .05 9.0

.04 .96 .43 23 .63 13 ..35 .20 .07 .03 .05 9.3

.08 1 .00 .42 23 ..76 12 .,67 .04 .08 .03 .05 4.2

.08 4.13 .61 20 ..21 9 ..92 1 .17 .04 .02 .02 11.6

.07 .99 .73 23 .,56 12 ..60 .21 .05 .03 .06 8.

.03 .74 .32 18 ..22 13 . 61 3 .53 .06 .03 . 04 4.2

.03 .76 .34 18 . 39 13 . 37 3 .70 .07 .03 .04 4.8

.04 1 .05 .42 24 . 02 13 . 64 .07 .09 .02 .03 9.0

.05 .98 .42 20 . 35 9 . 57 .18 .07 .03 .05 10.3

.04 1 .00 .35 19 .,17 1

. 86 2 .28 .06 .03 .05 9.8

.03 1 .55 .48 19 . 99 1

. 82 2 .33 .05 . 04 .06 11.1

.04 .91 .33 18 . 89 1

. 70 2 .30 .06 .03 .05 9.0

.05 1 .01 .44 20 . 29 9 . 77 .12 .06 .03 .05 10.6

.03 1 .05 .42 19 . 52 1

. 56 2 .45 .05 .04 .05 11.8

.07 1 .08 .81 23 . 86 12 . 61 . 1

.05 .03 .07 10.5

.03 2.25 .21 17 . 78 1

. 54 2 .29 .05 .03 . 04 6.5

.03 .90 .37 18 . 98 1

. 39 2 .21 .06 .03 .05 9.5

.05 1.01 .48 19 . 40 12 . 36 2 ,.41 .05 .04 .05 9.6

.07 .93 .35 23 . 50 12 . 94 .16 .08 .02 .04 8.5

.03 .99 .40 20 . 39 9 . 79 ,08 .05 .03 .05 10.8

.02 .97 .74 19 . 40 1

. 29 2 ..12 .06 .04 .05 10.3

.03 .96 .36 19 . 13 1

. 90 2 ..28 .05 .03 .05 10.6

.02 1.22 .36 19 . 78 1

. 70 2 .,21 .06 .03 .05 11.4

.03 .83 .32 18 . 85 11 . 35 2 . 12 .06 .03 .04 8.9

.03 1.21 .33 19 . 19 1

. 69 2 . 24 .08 .04 .05 9 . 1 .04 1 .22 .37 19 . 63 1

. 61 2 . 28 .08 .04 .04 9.3

.03 .95 .40 18 . 84 12 . 06 2 . 21 .06 .03 .05 8.7

271
Appendix I. FN Data Base, continued

C Mn Si Cr Ni N Mo Cb Ti FN

.03 .96 .36 19 .56 12 .00

.05 .98 .66 20 .62 10 .29 .03 .83 .39 20 ,16 9 .84

.03 1 .46 .37 20 .19 9 .82 .06 1 .56 .33 24 .22 12 .87

.03 .98 .73 20 . 13 10 .25

.03 .95 .68 20 .21 10 .26

.02 .89 .67 20 .01 10 .29 .03 1 .01 .65 20 .38 10 .49

.03 .92 .62 20 .01 10 .25

.02 .90 .68 20 .37 10 .21 .08 1 .00 .82 23 .82 12 .87

.06 .94 .68 23 .67 12 .79

.07 1 .07 .82 24 , 09 13 . 18

.05 .95 .43 19 .13 1

1

.61 .05 .86 .60 19 .87 9 .84

.05 .87 .58 20 .35 9 .96

.04 1 .22 .37 19 .63 1

1

.61 .04 1 .38 .39 19 .19 1

1

.70

.03 1.43 .38 19 .32 1

1

.57

.02 .92 .31 18 .85 1

1

.20 .03 .95 .33 19 .05 1

1

.35

.03 1 .00 .37 20 .02 10 .00

.03 .94 .36 20 .16 9 53

.05 .92 .35 20 .26 10 .03

.06 .94 .36 20 .06 9 .90

.05 1.05 .42 20 .38 9 .89 .04 .96 .41 20 .50 10 .22 03 .92 .39 20 .58 10 . 24

.02 1 .13 . 41 20 .20 9 .73

.03 .94 .31 20 , 04 10 .07

.02 1 .08 .38 20 .21 9 .71 .03 .. 1.04 .42 20 .,27 9 .91 .03 .86 .36 19 ..80 10 . 03

.04 1 .06 .45 20 . 31 9 .,76

04 1 .00 .43 20 . 32 9 . 97

03 1 .01 .40 20 . 06 9 70

03 1 .07 .44 20 . 35 9 . 92

03 1 .07 .43 20 . 53 9 95

03 1 .02 .38 20 . 32 10 . 07

07 .95 .46 24 . 25 12 . 75

07 .93 .41 23 . 60 12 . 47

07 .98 .43 23 . 99 12 . 64

08 .96 .35 23 . 95 13 . 22

03 1 .06 .44 22 . 78 13 . 62

07 .92 .40 23 . 49 12 . 39

07 .97 . 47 23 . 60 12 . 74

06 .96 .42 23 . 35 12 . 29

06 1 . 15 .38 20 . 28 9 . 71 05 .99 .37 20 . 21 9 . 49

03 1 .70 .53 20 . 88 17 . 45

03 1.15 .44 21 . 64 17 . 25

03 1 .26 .39 21 . 72 17 . 45

03 1 .29 .44 21 . 78 17 . 23

15 1 .42 .70 23 . 80 9 . 95

03 2.23 .31 18 . 12 12 . 25

02 2.31 .31 18 . 02 13 . 65

02 2.25 .25 18 . 19 13 . 60

02 2.04 .38 18 . 00 12 . 80

02 2.03 .22 18 . 10 12 . 80

02 2.31 .24 18 . 57 12 . 50

02 2.31 .24 18 . 05 12 . 50

02 2.33 .24 18 . 00 13 . 70

2.38 .08 .03 .05 8.7

. 1

1

.04 .03 .06 9.0

.16 .04 .03 .05 11.1

.12 .06 .02 .04 11.6

.09 .08 .02 .03 11.4

.17 .08 .03 .04 8.2

.17 .06 .03 .05 9.6

.15 .05 .03 .05 10.9

.20 .05 .03 .06 8.9

.20 .05 .03 .04 8.7

.16 .05 .03 .05 10.7

.12 .07 .03 .08 8.5

.20 .07 .02 .05 8.7

.12 .07 .02 .08 9.6

2.57 .06 .04 .05 8.5

.16 .05 .03 .04 8.4

.23 .05 .03 .04 9.4

2.28 .08 .04 .04 9.3

2.33 .06 .04 .05 8.8

2.23 .05 .04 . 04 12.1

2.33 .05 .04 .04 10.4

2.16 .08 .03 .04 8.2

.18 .06 .03 .05 9.1

. 18 .07 .03 .04 9.9

.22 .05 .03 .04 8.7

.16 .04 .03 .06 8.5

.14 .07 .03 .05 9.1

.17 .05 .03 .05 9.3

.17 .05 .03 .05 11.1

. 1

1

.08 .03 .04 11.6

.17 .06 .02 .04 8.7

. 1

1

.08 .03 .05 10.8

.25 .08 .03 .04 10.5

.17 .04 .03 .05 9.7

.20 .09 .03 .05 7.9

.20 .07 .03 .04 12.6

. 1

1

.08 .03 .04 11.3

.15 .08 .03 .05 8.0

.15 .09 .03 . 04 8.2

.17 .03 .03 .05 11.9

.24 . 1

1

.03 .05 8.5

.23 .08 .03 .06 9.9

.22 .08 .03 .06 10.6

.10 .05 .03 .05 8.3

.04 .06 .02 .07 6.6

.23 .08 .03 .05 9.6

.24 .09 .03 .06 8.8

.21 .07 .02 .05 10.1

.12 .08 .02 .05 6.9

.08 .08 .03 .05 9.3

3.82 .07 . 15 .03 .9 3.52 .07 .13 .03 .8 3.64 .07 .08 .03 .7 3.92 .07 .17 .03 .8 3.76 .12 .03 .04 24.7

2.05 .03 .03 .04 5.9

2.06 .03 .03 .04 1 .6 2.14 .03 .03 .04 .9 2.10 .04 .03 .04 2.2

2.10 .04 .03 .03 2.4

2.16 .02 .03 .04 3.4

2.16 .02 .03 .04 3.7

2.16 .03 .03 .04 1 .0

272
Appendix I . FN Data Base , continued

c Mn Si Cr Ni N Mo Cb Ti FN

.02 2 . 06 .31 18 . 15 12 . 30 2 .07 .03 .03 .03 5 . 8 .02 2 . 36 .25 18 . 07 13 . 75 2 ..04 .02 .03 .04 9 .04 12 ..40 .29 18 . 56 5 . 26 . 1

.16 .00 .00 9 . 8

.03 10 . 80 .51 18 . 56 4 . 88 .11 .25 .00 .00 6 . 2 .05 10 . 16 .36 17 . 69 5 . 00 .11 .19 .00 .00 4 . 9 .04 10 . 00 .38 18 . 16 4 . 61 . 1

.17 .00 .00 9 . 3

.04 1 . 35 .56 21 . 75 17 . 50 3 .80 .08 .14 .03 1 . 4 .03 1 . 30 .44 21 . 67 12 . 80 3 .82 .07 .02 .03 19 . 0 .03 .93 .44 22 . 76 12 . 90 2 .76 .06 .03 .01 19 . 0

.03 .95 .41 21 . 57 17 . 00 4 .00 .06 .01 .03 1 . 0 .03 90 .36 21 . 20 17 . 20 3 .84 .07 . 1

.03 1 . 0 .03 2 ..19 .42 17 . 85 12 . 30 2 .09 .04 .03 .04 4 . 7 .03 2 ., 17 .44 17 . 73 12 . 05 2 .20 .04 .04 .04 4 . 9 .04 2 .,09 .29 18 . 10 13 . 60 2 .04 .05 .04 .04 7

.03 2 ..21 .27 18 . 30 13 . 50 2 .02 .06 .03 .04 ,7 .03 2 ..37 .27 18 . 02 13 . 50 2 .07 .04 .03 .04 .7 .03 1 ..37 .27 17 . 66 12 ..60 2 .20 .04 .03 .05 3 ., 1 .03 1 ..40 .28 17 . 83 12 ..60 2 .18 .04 .04 .05 3 ..4 .03 2 .22 .32 17 . 71 12 . 09 2 .22 .04 .03 .04 5 ..4 .03 2 ..26 .30 17 . 74 12 . 63 2 .26 .03 .03 .04 3 ..6 .03 1 ,.64 .37 18 . 04 12 ..77 2 .31 .06 .04 .05 1 ..2 .04 1 .67 .37 18 . 41 12 ..75 2 .29 .10 .04 .05 6

.03 1 .64 .38 18 .,42 12 ..67 2 .33 .07 .04 .05 2 . . 1 .03 2 . 15 .20 18 ..28 1

..45 2 .23 .03 .03 .03 6 .6 .03 2 .05 .31 18 .,10 1

..60 2 .20 .03 .03 .04 6 . .4 .04 2 .16 .25 18 . 25 1

.21 2 .30 .04 .02 .04 6 .6 .04 2 . 18 .34 17 ..97 11 ,.80 2 .33 .05 .03 .04 4 .9 .05 4 .70 .50 19 ..54 13 ,.10 4 .77 . 19 .09 .01 .9 .04 2 .22 .28 17 .95 13 .10 2 .21 .04 .03 .04 .9 .03 1 . 67 .42 18 ,, 45 12 .60 2 .16 .06 .04 .06 4 .2 .03 1 .72 .52 19 .65 12 . 10 2 . 18 .06 .04 .07 8 .8 .03 2 .21 .28 17 .92 13 .09 2 . 1

.04 .03 .04 1 . 0 .04 - 2 .25 .30 18 .01 13 . 17 2 .12 . 04 .03 .04 1 . 1 .15 1 .42 .49 20 ,.45 10 .00 . 1 1 .02 .03 .08 2 . 6 .05 1

.23 .39 17 .89 4 .82 .17 .19 .02 .02 5 .4 .04 1

.49 .28 17 ,.81 5 .14 .18 .18 .02 .02 4 .9 .04 1 .45 .28 17 .69 12 .25 2 .13 .08 .03 .04 2 .0 .04 1 .51 .36 17 .46 12 .51 2 .17 .05 .04 .05 1 .0 .06 1 .59 .41 18 .74 12 . 10 2 .27 .08 .05 .06 3 .3 .03 1 .10 .36 23 .35 13 .25 .09 .05 .03 .04 8 . 1 .05 1 .29 .34 19 .69 9 .43 .25 .05 .05 .04 9 .8 .05 1 .36 .40 19 .70 9 .84 .25 .05 .05 .04 9 . 0 .05 1 .33 .37 19 .64 9 .75 .25 .05 .05 .04 8 .5 .05 1 .35 .38 19 .66 9 .65 .25 .06 .05 .04 7 .3 .12 1 .70 .46 19 .82 9 .56 1 .98 .04 .02 .04 1

.3 .07 1 .64 .53 23 .97 13 .09 .16 .07 .02 .05 7 .5 .03 1 .08 .46 23 .19 13 .55 .10 .06 .03 .07 7 . 4 .04 1 .05 .45 23 .08 13 .38 .09 .06 .03 .06 6 .3 .05 .98 .36 20 .16 9 .30 .08 .04 .03 .06 1 1 .2 .06 1 .27 .42 20 .15 9 .43 .14 .07 .03 .06 8 .3 .06 1 .15 .38 20 .28 9 .71 .12 .08 .02 .05 6 .9 .02 .90 .68 20 .37 10 .21 .16 .05 .03 .05 10 . 7 .07 1 .02 .87 23 .95 12 .94 .13 .07 .03 .06 8 .8 .05 1 .03 .65 20 . 1

10 .17 .16 .04 .03 .05 1 1 . 3 .03 .79 .32 18 .77 13 .44 3 .40 .07 .03 .02 4 .8 .06 .95 .61 19 .77 10 .00 .12 .06 .03 .05 6 .3 .03 1 .45 .34 17 .38 12 .31 2 .22 .04 .04 . 04 3 .3 .08 .92 .36 23 .76 13 .30 . 1

.05 .03 . 06 8 . 7 .04 1 .17 .49 23 .63 13 .41 .12 .07 .03 .06 8 .8 .06 1 .03 .46 20 .32 10 .08 .16 .06 .03 .05 8 .8 .06 1 .34 .42 19 .58 9 .83 . 1

.06 .02 .04 7 .3 .05 1 .25 .47 19 .66 10 .12 . 1

.08 .02 .04 5 .5 .03 1 .40 .38 18 .31 13 .13 2 .29 .06 .03 .03 1 .3

273
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.04 1 .12 .43 19 .59 9 .84 . 1

1

.07 .54 .05 6 .3 .07 .91 .83 23 .39 13 . 1 1 .18 .08 .03 .09 6 .8 .05 1 .05 .41 19 .87 9 ,81 .08 .05 .57 .05 9 0 .06 1 .12 .36 19 .63 9 .52 .31 .06 .03 .05 6 .0 .07 1 .20 .57 20 .00 9 .59 .32 .05 .04 .06 8 .2 .06 1 .35 .77 20 . 40 9 .51 .32 .06 .04 .07 10 .9 .04 .92 .45 19 .02 9 .47 . 1

1

.06 .55 .05 7 .2 .04 1 .15 .56 19 .80 9 .45 .12 .05 .58 .06 10 .2 .06 1 .47 .53 20 .18 9 .64 . 1

1

.08 .67 .04 8 .8 .05 1 .75 .46 19 .78 1

1

.50 .18 .05 .03 .04 3 .0 .05 1 .67 .42 19 .81 1

1

.45 1 .86 .05 .03 .04 7 .3 .05 1 .65 .41 19 .78 11 .60 2 .59 .05 .03 .04 9 . 6 .06 1 .55 .46 19 .81 1

1

.60 3 .68 .05 .03 .04 12 0 .06 1 .53 .46 19 .77 1

1

.55 5 .29 .05 .03 .04 15 .8 .06 1 .45 .41 19 .52 1

1

.60 6 .85 .05 .03 .04 20 .5 .03 1 .57 .29 17 .58 9 .19 .02 .05 .00 .00 2 .8 .03 1 .49 .34 15 .37 9 .06 2 .03 .03 .00 .00 3 .4 .03 3 .23 .39 17 .89 9 .02 .02 .04 .00 .00 4 0 .04 6 .54 .37 17 . 61 9 .13 .02 .05 .00 .00 2 .8 .04 6 .31 .34 15 .39 9 . 14 1 .99 .05 .00 .00 2 .8 .04 6 .31 .34 15 .39 9 .14 1 .99 .05 .00 .00 2 .8 .04 6 .27 .34 15 . 17 9 . 14 3 .84 .16 .00 .00 2 .6 .03 1 .66 .32 18 .36 8 .94 .02 .10 .00 .00 1 .8 .03 2 .96 .36 18 .13 9 .12 .02 .04 . 00 .00 3 .6 .04 6 .46 .35 18 .51 9 .17 .02 .10 .00 .00 2 .6 .04 9 .52 .38 17 .71 9 ,.10 .02 . 1

1

.00 .00 1 5 .03 3 .18 .34 18 .99 9 .05 .02 .15 .00 .00 1 4 .04 6 ,.62 .38 19 .51 9 06 .02 .15 .00 .00 2 2 .04 6 .61 .32 17 .77 9 . 10 1 .66 .17 .00 .00 2 .5 .04 9 .38 .35 18 .98 9 . 14 .02 . 16 .00 .00 1 . 7 .03 6 .. 1 1 .36 20 . 90 9 .29 .02 .25 .00 .00 1 . 6 .04 9 .56 .37 20 .22 9 .07 .02 .27 .00 .00 1 .2 .03 - 2 .90 .30 18 .79 9 .12 .02 .10 .00 .00 2 .0 .03 3 ,.25 .39 17 .58 9 .00 .02 .03 .00 .00 3 .2 .03 3 .16 .42 17 .90 9 .05 .02 .04 .00 .00 3 .0 .04 9 .55 .34 17 .70 9 ,.26 .02 .04 .00 .00 3 .6 .03 3 .10 .36 17 .75 9 .07 .02 .15 .00 .00 .8 .03 3 .46 .37 17 .82 9 . 10 .02 .15 .00 .00 .6 .03 9 .72 .34 17 .60 9 06 .02 .26 .00 . 00 2 .04 1 ,.29 .26 21 .02 5 37 3 . 00 .15 .00 .15 49 .0 .02 1 .63 .26 20 .99 6 .75 3 .09 .10 .00 .13 51 ,0

.03 1 ,.80 .30 20 .80 7 ..45 2 .92 .09 .00 . 13 42 0

.03 1 .70 .28 20 .33 7 ,90 2 .80 .09 .00 .13 27 0 .03 1 .95 .32 22 .23 7 .96 3 . 40 .12 .00 .13 43 0 .03 1 .98 .32 22 .24 8 . 40 3 .37 . 1

1

.00 .08 35 . 0 .02 1 .60 .42 22 .07 8 .70 3 .88 .13 .00 .23 43 . 0 .02 1 . 80 .61 21 .92 8 . 55 3 .26 .15 .00 .23 35 . 0

.02 1 75 .62 22 .29 8 . 55 3 .50 .13 .00 .33 62 . 0 .03 1 .90 .62 21 .85 8 . 55 3 . 40 . 14 .00 .54 83 . 0 .02 1 ..60 .44 21 .93 8 . 35 3 . 40 .15 .00 .13 32 . 0 .02 1 ..59 .52 22 .14 8 . 45 3 .42 .17 .00 . 13 37 . 0 .02 1 .49 .38 22 .31 8 ..50 3 .48 .17 .00 .05 27 . 0 .02 1 .,38 .39 21 .93 8 . 40 3 .32 . 17 .00 .05 24 . 0 .04 1 ..33 .47 24 .41 10 . 60 3 .28 . 1

1

.00 .21 43 . 0 .05 1 ..12 .56 24 .71 9.,10 3 .32 .16 .00 . 1

1

35 . 0

.05 1 . 14 .56 24 .24 9. 45 4 .00 .17 .00 . 1

1

37 . 0 .05 .95 .42 24 .62 10 . 80 4 .08 .17 .00 . 1

1

28 . 0 .02 .90 .21 25 .28 9 20 4 .27 .12 .00 . 13 56 . 0

.03 91 .17 25 .22 9. 35 4 .27 . 1 1 .00 .01 67 . 0

.02 1 ..00 .25 25 .25 9 . 30 3 .46 . 1

1

.00 .13 57 . 0 .02 97 .32 24 .79 10 . 25 3 .64 . 14 .00 .02 34 . 0 .02 .77 .31 24 ,.58 10 . 95 3 .56 .12 .00 .02 27 . 0

.01 1 . 06 .36 24 .51 10 . 17 3 . 18 .08 .00 .02 40 . 0

274
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.06 1 .53 .40 20 .57 9.29 .05 .05 .02 .02 7 .0 .06 1.44 .39 20 .57 9.31 .05 .06 .02 .02 7 .0 .06 1 .51 .40 20 .48 9.35 .05 .07 .02 .02 7 .0 .07 1 .45 .39 20 .19 9.41 .04 .06 .02 .02 6 .0 .06 1 .40 .39 20 .09 9.38 .04 .06 .02 .02 4 .0 .07 1 .49 .40 20 .38 9.41 .04 .09 .02 .02 5 .0 .06 1 .48 .41 20 .19 9.33 .04 .09 .02 .02 5 .0 .06 1 .45 .42 20 .47 9.34 .05 .07 .02 .02 8 .0 .06 1 .35 .31 20 .20 9.37 .03 .06 .01 .02 5 .0 .06 1 .48 .39 20 .31 9.32 .03 .07 .01 .02 6 .0 .06 1 .44 .36 20 .37 9.37 .03 .07 .01 .02 6 .0 .06 1 .58 .42 20 .73 9.26 .03 .06 .02 .02 8 .0 .06 1.33 .29 20 .27 9.41 .03 .09 .01 .01 6 .0 .06 1 .49 .39 20 .36 9.37 .03 .08 .02 .02 6 .0 .06 1 .55 .36 20 .42 9.45 .12 .10 .01 .02 7 .0 .06 1 .60 .37 20 .58 9.35 .03 . 10 .01 .02 7 .0 .06 1 .49 .37 20 .64 9.39 .04 .08 .02 .02 10 .0 .06 1 .50 .37 20 .64 9.34 .03 .08 .02 .02 9 .0 .06 1 .51 .38 20 .63 9.81 .12 .08 .02 .02 10 .0 .06 1 .53 .40 20 .49 9.76 .12 .08 .02 .02 9 .0 .06 1 .50 .40 20 .33 9.72 .12 .06 .02 .03 7 .0 .06 1 .47 .38 20 .31 9.75 .12 .09 .02 .02 6 .0 .06 1 .51 .42 20 .41 9.71 .12 .05 .02 .02 7 .0 .06 1 .50 .40 20 .33 9.72 .12 .06 .02 .03 7 .0 .06 1 .50 .41 20 .39 9.70 .12 .05 .02 .02 6 .0 .05 1 .51 .43 20 .41 9.73 .12 .05 .02 .02 8 .0 .07 1.38 .41 20 .16 9.32 .03 .08 .02 .02 5 .0 .06 1.42 .38 20 .17 9.45 .03 .06 .02 .02 5 .0 .06 1 .49 .43 20 .32 9.40 .03 .07 .02 .02 5 .0 .07 1 .53 .43 20 .25 9.39 .04 .07 .02 .03 5 .0 .07 1.52 .42 20 .28 9.31 .04 .06 .02 .02 6 .0 .07 1 .48 .40 20 .26 9.33 .05 .07 .02 .02 5 .0 .06 1 .52 .45 20 ..48 9.32 .05 .07 .02 .02 8 .0 .06 1 .47 .42 20 .42 9.39 .05 .08 .02 .02 6 .0 .06 1.47 .43 20 ..57 9.31 .05 .07 .02 .02 8 .0 .03 1.69 .46 23 ..14 13.44 .02 .04 .02 .03 9 .0 .03 1 .71 .46 23 ..00 13.41 .02 .05 .02 .03 9 .0 .03 1 .69 .48 23 .,20 13.57 .02 .07 .02 .03 8 ,0

.04 1 .69 .43 22 ..94 13.64 .02 .07 .02 .02 6 .0 .03 1 .70 .47 22 ..94 13.78 .06 .07 .02 .00 7 .0 .03 1 .71 .45 22 . 93 13.73 .06 .07 .02 .03 5 .0 .02 1 .57 .42 23 ..10 13.85 .05 .09 .02 .02 8 ..0 .06 1 .46 .44 20 . 43 9.33 .03 .08 .02 .03 8 .,0 .06 1 .48 .44 20 . 35 9.45 .03 .08 .02 .02 4 . 0 .06 1 .55 .43 20 . 34 9.42 .02 .10 .03 .03 4 . 0 .06 1 .40 .42 20 . 22 9.40 .03 .13 .02 .02 4 . 0 .07 1 .67 .50 20 . 62 9.28 .02 . 10 .03 .03 7 . 0 .07 1 .68 .44 20 . 76 9.31 .02 .09 .03 .03 5 . 0 .06 1 .54 .48 20 . 41 9.18 .02 .09 .02 .03 6 . 0

.06 1.48 .46 20 . 64 9.46 .04 . 1

.03 .05 4 . 0 .06 1.33 .40 20 . 29 9.33 1 .50 .09 .20 .20 5 . 0 .03 1 .80 .48 23 . 24 13.88 .09 .05 .02 .03 12 . 0

.03 1.79 .48 23 . 50 13.50 .09 .04 .02 .05 10 . 0 .03 1 .79 .50 22 . 89 13.70 .09 .04 .02 .04 8 . 0 .03 1 .84 .50 23 . 43 13.49 .09 .05 .03 .05 9 . 0

.03 1 .75 .40 22 . 79 13.75 .04 .05 .02 .03 7 . 0 .03 1.56 .41 22 . 83 13.58 .03 .07 .03 .02 7 . 0 .03 1 .72 .44 22 . 98 13.64 .04 .05 .02 .03 8 . 0

.03 1 .73 .49 23 . 84 13.61 .05 .06 .02 .03 1 1 . 0 .03 1 .76 .47 23 . 12 13.49 .06 .06 .02 .02 7 . 0 .03 1 .73 .44 23 . 23 13.39 .06 .06 .02 .02 1 1 . 0

.03 1 .69 .44 23 . 38 13.38 .06 .04 .02 .03 10 . 0 .04 1 .71 .44 23 . 24 13.33 .06 .07 .02 .03 10 . 0

275
Appendix I. FN Data Base, continued

C Mn Si Cr Ni N Mo Cb Ti FN

.03 1 .57 .30 23 .46 13 .78 .03 1 .51 .44 24 .28 12 .71 .04 1.78 .54 22 .42 13 .75

.03 1 .85 .53 23 .18 13 .33

.03 1 .80 .50 22 .91 13 .45

.03 1 .68 .46 22 .95 13 .45

.03 1 .75 .50 23 .03 13 .43 .03 1 .79 .44 22 .61 12 .75

.03 1 .97 .55 23 .00 12 .72

.04 1 .79 .44 22 .73 12 .75

.07 1 .76 .39 23 . 44 12 .61 .07 1.78 .42 23 .34 12 .66

.07 1 .83 .41 23 .61 12 .94

.07 1 .81 .48 23 .33 12 .90

.07 1 .79 .39 23 .74 12 .98

.07 1 .46 .41 23 .06 13 .02

.07 1 .34 .31 23 .01 13 .25

.07 1 .34 .33 23 .05 12 .99 .03 1 .63 .40 23 . 16 13 .34

.03 1 .57 .38 23 . 1

1

13 .33 .07 1.43 .39 23 .26 13 . 14

.07 1 .29 .34 23 .05 12 .88

.08 1.43 .39 23 .21 13 .12 .07 1 .33 .39 23 .23 12 .90

.07 1 .39 .40 23 .27 12 .87

.07 1 .26 .33 23 .22 12 .91 .06 1 .37 .40 23 .32 12 .87 .08 1 .97 .44 23 35 12 . 68

.07 1 .73 .41 23 .90 12 .61 .07 1 .75 .42 23 .93 12 .52 .07 1 .71 .41 23 .77 12 .57 .07 1 .98 .44 23 .63 12 .56 .07 " 1.79 .46 24 ,.21 12 ,.52 .07 1 .83 .48 24 ,36 12 .54 07 1 .29 .36 22 . 96 12 . 92

08 1 .50 .46 23 08 13 . 04

.08 1 .55 .59 23 . 28 13 . 08

.07 1 .38 .38 23 . 16 13 . 19 .07 1 .36 .38 22 . 84 12 . 90

.07 1 .61 .48 23 . 54 12 . 90

07 1 .61 .47 23 . 33 12 . 77

06 1 .47 . 40 23 . 02 12 . 50

.06 1 .63 . 47 23 . 22 12 . 51 07 1 .50 .43 22 . 86 12 . 48

07 1 .58 .51 23 . 12 12 28

06 1 .66 .49 23 . 36 12 25

07 1 .00 .52 23 . 28 12 . 29

06 1 .66 .52 23 . 56 12 . 25

06 1 .50 .46 23 . 33 12 . 30

06 1 .72 .52 23 . 60 12 . 16 06 1 .63 . 40 23 . 58 12 . 88

06 1 .63 .42 23 . 19 12 . 85

06 1 .54 .41 23 . 12 12 . 87

06 1 .61 .43 23 . 17 12 . 87

06 1 .66 .44 23 . 54 12 . 82

06 1.70 .43 23 . 38 12 . 80

06 1 .71 .43 23 . 50 12 . 81 07 1 .60 .44 23 . 36 13 . 00

07 1 .55 .42 23 . 30 13 . 09

07 1 .58 .46 23 . 24 12 . 96

07 1 .87 .44 23 . 51 12 . 69

07 1 .88 .44 23 . 89 12 . 46

08 1 .92 .47 23 . 86 12 . 49

. 10 .06 .02 .02 7.0

.04 .05 .02 .03 20.0

.03 .07 .02 .03 4.0

.02 .05 .03 .03 9.0

.02 .06 .02 .03 8.0

.02 .09 .02 .03 8.0

.02 .06 .02 .03 8.0

2.51 .04 .02 .02 12.0

2.27 .06 .02 .04 7.0

2.82 .06 .02 .03 5.0

.02 .07 .02 .02 8 0 .02 .07 .02 .03 7.0

.02 .07 .02 .02 8.0

.02 .08 .02 .03 5.0

.01 .07 .02 .03 5.0

.03 .09 .02 .02 3.0

.03 .08 .02 .02 5.0

.03 .08 .02 .02 5.0

.06 .06 .02 .03 11.0

.06 .07 .02 .02 9.0

.03 .10 .20 .20 4.0

.03 .07 .02 .02 6.0

.03 .08 .03 .02 6.0

.03 .07 .02 .02 6.0

.03 .07 .02 .02 6.0

.03 .08 .02 .02 8.0

.03 .06 .02 . 02 6.0

.03 .08 .02 .03 5.0

.03 .06 .02 .03 8.0

.03 .07 .02 .03 7.0

.02 .07 .02 .03 7.0

.03 .08 .02 .03 7.0

.01 .07 . 02 .03 9.0

.01 .08 .02 .03 9 0 .07 .08 .02 .02 6.0

.03 .07 .02 .02 3.0

.03 .07 .02 .03 4 . 0 .03 . 06 .02 .02 4.0

.02 . 1

1

.02 .02 3.0

.03 07 .02 .03 6.0

.06 .07 .02 .03 6.0

.02 .08 .02 .02 6.0

.02 .08 .02 .02 5.0

.02 .07 .02 .02 5.0

.03 .06 .01 .02 8.0

.03 .06 .02 .02 8.0

.03 .10 .02 .02 7.0

.03 .07 .02 .03 8.0

.03 .07 .02 .02 7.0

.03 .05 .02 .03 8.0

.02 . 10 .01 .02 6.0

.02 .09 .01 .02 5.0

.02 .09 .02 .02 5.0

.02 .07 .01 .02 5.0

.02 .07 .02 .03 6.0

.03 .08 .02 .03 5.0

.03 .08 .02 .03 5.0

.03 .07 .02 .02 5.0

.03 .07 .02 .02 4.0

.03 .07 .02 .02 4.0

.02 .07 .02 .02 6.0

.02 .07 .02 .02 8.0

.02 .08 .02 .03 7.0

276
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.08 1.93 .47 23.93 12.41 .02 .07 .02 .03 7.0

.07 1 .96 .46 24.15 12.72 .02 .09 .02 .03 8.0

.07 2.09 .50 24.26 12.49 .01 .07 .02 .03 8.0

.07 2.14 .51 24.32 12.47 .01 .07 .02 .03 8.0

.07 2.12 .56 24.52 12.69 .02 .08 .03 .03 8.0

.08 1.75 .43 24.17 12.45 .02 .08 .02 .03 7.0

.08 1.75 .40 24.15 12.69 .02 .08 .02 .03 6.0

.08 1 .76 .43 24.14 12.57 .01 .08 .02 .03 7.0

.08 1 .79 .44 24.28 12.63 .02 .08 .02 .03 8.0

.07 1 .76 .46 24.27 12.44 .02 .04 .02 .03 7.0

.07 1 .59 .45 23.12 12.68 .02 .09 .02 .02 8.0

.07 1 .64 .46 23.17 12.65 .02 .08 .02 .02 5.0

.07 1.68 .49 23.23 12.67 .02 .08 .02 .03 5.0

.07 1 .67 .45 22.93 12.72 .02 .07 .02 .02 4.0

.06 1 .55 .44 23.00 12.63 .02 .08 .02 .02 4.0

.06 1 .55 .46 23.00 12.57 .03 .09 .02 .03 5.0

.03 1 .65 .43 18.64 11.33 2.29 .04 .02 .02 9.0

.02 1 .62 .41 18.65 11.41 2.28 .06 .02 .02 7.0

.02 1 .63 .43 18.80 11.37 2.30 .06 .02 .02 8.0

.02 1 .69 .38 18.80 11.62 2.24 .06 .01 .02 7.0

.02 1 .68 .35 18.68 11.70 2.25 .07 .01 .02 7.0

.02 1 .32 .31 18.62 11.37 2.20 .06 .01 .01 6.0

.02 1 .46 .38 18.83 11.37 2.18 .04 .01 .02 9.0

.03 1 .52 .36 19.17 11.35 2.15 .05 .02 .02 8.0

.02 1 .42 .34 18.94 11.42 2.19 .06 .01 .02 7.0

.02 1 .46 .37 18.50 11.39 2.18 .06 .01 .02 8.0

.02 1 .43 .35 18.86 11.42 2.18 .05 .01 .02 8.0

.06 1 .70 .48 19.27 11.39 2.32 .06 .02 .02 7.0

.06 1 .61 .43 19.47 11.59 2.28 .07 .02 .02 7.0

.06 1 .70 .47 19.53 11.35 2.32 .05 .02 .02 9.0

.06 1 .64 .46 19.47 11.47 2.33 .06 .02 .02 5.0

.06 1 .57 .43 19.14 11.64 2.26 .07 .02 .02 4.0

.02 -1.58 .42 18.90 11.48 2.17 .07 .02 .02 8.0

.02 1.57 .38 19.01 11.41 2.17 .06 .02 .02 8.0

.02 1 .65 .42 19.02 11.30 2.15 .06 .01 .02 8.0

.02 1 .49 .39 19.00 11.43 2.17 .07 .02 .02 6.0

.03 1 .32 .33 18.76 11.60 2.15 .07 .02 .01 5.0

.02 1 .51 .42 18.73 11.39 2.15 .05 .02 .02 6.0

.02 1 .52 .40 19.10 11.54 2.15 .06 .02 .02 7.0

.02 1 .80 .45 18.91 11.51 2.27 .05 .02 .02 9.0

.02 1 .76 .49 19.02 11.34 2.27 .05 .02 .02 9.0

.02 1 .69 .43 18.89 11.46 2.28 .05 .02 .02 9.0

.02 1 .68 .45 18.92 11.42 2.31 .06 .02 .02 7.0

.02 1 .65 .46 18.75 11.19 2.32 .05 .02 .02 9 . 0 .06 1 .42 .46 19.38 11.41 2.26 .07 .01 .01 5.0

.05 1 .39 .46 19.62 11.25 2.32 .09 .01 .02 7.0

.07 1 .59 .46 19.56 11.39 2.25 .06 .01 .02 5.0

.07 1.51 .46 19.34 11.43 2.25 .06 .02 .02 5.0

.06 1.50 .45 19.48 11.43 2.26 .07 .01 .02 6.0

.06 1 .40 .37 19.66 11.49 2.26 .07 .01 .01 5.0

.03 1.79 .36 19.27 12.35 3.56 .08 .01 .01 8.0

.02 1 .87 .40 19.73 12.49 3.44 .05 .01 .01 6.0

.03 1 .93 .45 19.82 12.38 3.42 .06 .01 .02 9.0

.03 1 .87 .40 19.37 12.29 3.43 .05 .01 .01 8.0

.03 1 .86 .40 19.42 12.42 3.42 .07 .01 .02 8.0

.05 1 .80 .44 19.28 11.46 2.18 .05 .01 . 01 4 . 0 .06 1 .64 .43 19.48 11.37 2.27 .07 .02 .03 4 . 0 .06 1 .50 .40 19.37 11.53 2.29 .07 .02 .02 5.0

.06 1.48 .41 19.34 11.40 2.30 .08 .02 .02 5.0

.06 1 .69 .47 19.72 11.19 2.26 .08 .02 .03 6.0

.06 1.67 .44 19.37 11.55 2.29 .06 .02 .02 5.0

.07 1.63 .44 19.38 11.56 2.30 .07 .02 .02 4.0

.06 1 .62 .42 19.59 11.53 2.30 .07 .02 .03 5.0

277
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.03 1 .63 .39 19.97 9.74 .08 .05 .54 .02 10.0

.03 1.50 .34 19.95 9.86 .08 .12 .60 .20 11 .0 .03 1 .66 .36 19.65 9.93 .06 .06 .52 .02 7.0

.03 1 .55 .34 19.75 10.05 .05 .06 .54 .02 8.0

.04 1 .82 .41 19.52 12.46 3.67 .07 .01 .01 7.0

.03 1 .84 .43 19.74 12.48 3.69 .05 .01 .02 8.0

.03 2.24 .50 20.06 12.41 3.56 .05 .01 .02 8.0

.02 2.00 .44 19.71 12.52 3.60 .05 .01 .02 8.0

.02 2.05 .48 19.95 12.31 3 60 .07 .02 .02 2.0

.02 1 .94 .44 19.66 12.49 3.68 .06 .02 .02 7.0

.02 2.03 .42 19.70 12.40 3.68 .07 .01 .02 8.0

.03 2.02 .44 19.54 12.51 3.57 .07 .01 .02 7.0

.03 1 .66 .38 19.61 13.23 3.91 .10 .01 .02 2.0

.03 1 .89 .39 20.15 12.40 4.34 .06 .01 .02 3.0

.03 1 .60 .42 19.58 9.79 .04 .09 .50 .02 7.0

.02 1 .63 .41 19.84 9.89 .06 .06 .46 .02 1 .0 .03 1 .65 .49 20.18 9.91 .03 .05 .59 .02 1 .0 .02 1 .84 .21 20.20 10.25 .03 .05 .71 .02 9.0

.02 1 .61 .36 19.74 10.20 .05 .09 .59 .02 8.0

.02 1 .80 .40 19.94 10.18 .05 .08 .64 .02 6.0

.02 1 .73 .39 19.87 10.22 .05 .09 .63 .02 8.0

.02 1 .76 .42 19.71 10.19 .05 .08 .63 .02 8.0

.02 1 .90 .43 20.01 10.12 .05 .07 .70 .03 10.0

.04 1 .89 .68 23.21 15.27 3.48 . 10 .00 .00 6.0

.04 1 .85 .70 23.07 14.29 3.97 .10 .00 .00 6.8

.04 1.87 .73 24.61 14.34 4.51 .09 .00 .00 5.6

.04 1 .82 .63 24.54 13.85 4.03 .09 .00 .00 3.6

.03 1 .45 .62 22.48 15.78 3.97 .08 .00 .00 2 4 .02 1 .22 .39 22.57 14.97 2.73 .13 .00 .00 5.2

.03 1 .35 .64 21 .94 16.38 5.31 .08 .00 .00 .6 .02 1 .20 .94 22.10 15.52 3.38 .12 .00 .00 1 .0

. 04 1 .62 .65 24.37 16.31 5.02 . 10 .00 .00 .0 .03 1 .64 .68 24.74 14.02 4.32 .09 .00 . 00 7.4

.04 1 .66 .64 23.39 14.46 4.73 .09 .00 .00 2.0

.03 1 .71 .64 23.08 14.41 4.06 .09 .00 .00 6.0

.04 1 . 69 .73 23.13 13.73 3.20 . 1

1

.00 .00 17 0 .04 1 . 68 .68 23.44 13.64 3.23 . 10 .00 .00 18.2

.03 1 .70 .69 23.09 13.93 3.10 . 10 .00 .00 13.7

.03 1 .65 .72 24.73 14.43 4.53 .09 .00 .00 9.2

.03 1 . 64 .68 24.74 14.02 4.32 .09 .00 .00 7.4

.04 2.00 .56 19.85 9.77 .28 .04 .00 .00 8.0

.05 1 .91 .51 19.68 9.30 .16 .06 .00 .00 7.8

.05 1 .85 .49 19.95 9 . 40 . 19 .05 .00 .00 7.6

.03 2.13 .55 24.30 13.41 .22 .09 .00 .00 9 4 .02 1 .87 .54 19.98 9.70 .16 .06 .00 .00 10.0

.02 1 .89 .54 20.03 9.62 .16 .06 .00 .00 9 . 4 .02 1 .87 .55 20.05 9.62 .16 .06 .00 .00 9 . 6 .04 1 .99 .57 19.84 9.77 .29 .04 .00 . 00 8.7

.03 1 .87 .53 18.80 12.28 2.11 .05 .00 .00 5.5

.02 1 .79 .53 20.51 9.80 .16 .07 .00 .00 10.0

.03 1 .91 .68 18.11 12.46 2.14 .05 .00 .00 3.9

.03 1 .95 .60 19.51 9.60 .15 .05 .00 .00 7.8

.10 1 .85 .61 24.49 12.89 .19 . 10 .00 .00 6.8

.06 1 .88 .57 19.76 9.34 .19 .08 .00 .00 6.8

.09 2.04 .55 24.40 12.85 .17 .12 .00 .00 7.8

.10 2.16 .57 24.58 12.78 .17 .10 .00 .00 6.6

.05 1 .94 .50 20.13 9.26 .18 .06 .00 .00 9.8

.05 1 .93 .52 20.01 9.34 . 18 .07 .00 .00 8.4

.09 2.07 .56 24.31 12.87 . 17 . 1 1 .00 .00 6.0

.05 1 .93 .52 20.36 9. 17 .19 .06 .00 .00 9.2

.03 2.08 .59 20.27 9.42 .08 .07 .00 . 00 8.5

.03 1 .95 .53 20.12 9.18 .30 .08 .00 .00 9.2

.03 1 .85 .52 20.04 9.27 .30 .08 .00 .00 9.2

278
Appendix I. FN Data Base, continued

C Mn Si Cr Ni N Mo Cb Ti FN

.03 1 .84 .53 19 .99 9 .33

.10 2.11 .63 24 .53 12 .83

.03 1.79 .50 19 .90 9 .47

.07 2.04 .53 20 .04 9 .35

.06 2.01 .53 19 .89 9 .38

.08 1.22 .84 24 .97 12 .86

.04 1.86 .54 19 .71 9 .53

.05 1 .92 .53 19 .86 9 .37

.04 1.76 .52 19 .57 9 .40

.06 1.82 .49 19 .80 9 .29

.06 1 .78 .49 19 .72 9 .33

.06 1 .78 .50 19 .73 9 .39

.10 1 .92 .54 24 .58 13 .12

.09 1 .94 .56 24 .36 13 .04

.06 1 .75 .55 19 .79 9 .44

.06 1.73 .56 19 .71 9 .32

.09 1 .82 .51 24 . 19 13 .08

.05 1 .83 .54 20 .35 9 .29

.06 1 .90 .61 19 .77 9 .29

.06 1.92 .54 20 .01 9 .34

.06 1 .84 .51 19 .82 9 .29

.08 1 .83 .72 24 .66 13 . 13

.05 1.82 .52 20 .41 9 .20

.08 1 .50 .70 23 .02 12 .65

.03 2.01 .76 20 .60 9 .50

.09 2.22 .78 23 .83 12 .83

.09 2.09 .67 24 .58 12 .96

.09 1 .93 .52 24 .16 12 .98

.04 1 .23 .84 20 .39 9 .27

.05 1 .62 .45 18 .84 13 .05

.05 1 .70 .46 1

,.73 13 .00

.03 -1 .91 .61 18 .15 13 ,.54

.03 1 .82 .50 20 ,.21 9 ,.58 .03 1 .61 .46 18 ..30 12 ..55 .03 1 .67 .43 18 ..97 12 ..28

.03 2.02 .89 20 . 04 9 .,75

.03 1 .82 .50 20 . 21 9 . 58

03 1 .32 .21 17 . 33 12 . 37

04 1 .54 .22 17 . 58 12 . 47

03 1 .78 .22 17 . 73 12 . 82

03 2.33 .35 17 . 75 12 . 78

03 2.43 .22 17 . 79 12 . 96

03 2.25 .16 17 . 65 12 . 75

03 1 .43 .26 17 . 40 12 . 74

03 1 .56 .42 17 . 48 12 . 59

02 1 .63 .46 17 . 46 13 . 04

03 1 .56 .43 17 . 54 13 . 78

03 1 .74 .46 17 . 42 12 . 20

03 1 .65 .23 17 . 34 12 . 42

03 1 .70 . 15 17 . 30 12 . 54

03 1 .62 .51 17 . 48 14 . 31 03 1 .67 .61 17 . 34 14 . 84

03 1 .60 .35 17 . 42 14 . 45

03 1.66 .40 17 . 44 14 . 89

03 1 .56 .50 17 . 54 14 . 41 03 1 .56 .47 18 . 21 13 . 80

03 1 .68 .44 18 . 49 15 . 16 03 1 .75 .43 18 . 03 14 . 20

03 1 .63 .44 17 . 97 13 . 68

02 .71 .78 18 . 81 9 . 92

02 .35 .84 18 . 21 10 . 15

02 .37 .84 18 . 57 10 . 16

02 .37 .84 18 . 37 10 . 1

.30 .08 .00 .00 10 .6 .17 .12 .00 .00 6 .8 .30 .10 .00 .00 9 .8 .30 .07 .00 .00 7 .0 .30 .07 .00 .00 8 .2 .18 .12 .00 .00 10 . 1 .60 .05 .00 .00 6 .6 .30 .07 .00 .00 8 .0 .60 .05 .00 .00 10 .4 .16 .06 .00 .00 7 .4 .16 .06 .00 .00 7 .4 .16 .07 .00 .00 7 .8 .10 .06 .00 .00 8 .6 .20 .07 .00 .00 8 .2

.21 .07 .00 .00 5 .6 .20 .07 .00 .00 7 .3 .20 .07 .00 .00 7 . 0 .16 .06 .00 .00 10 .9 .16 .05 .00 .00 8 .0 .29 .05 .00 .00 7 .2 .16 .08 .00 .00 7 .0 .12 .07 .00 .00 9 .0 .16 .05 .00 .00 1

.6 2 .20 .06 .00 .00 13 .4 .14 .08 .88 .00 9 ,8

.08 . 1

.84 .00 9 .2 .14 .09 .00 .00 6 .8 .21 .07 .00 .00 6 . 4 .16 .07 .00 .00 12 .5

2 .13 .08 .00 .00 2 .4 2 . 14 .07 .00 .00 1 .5 2 . 10 .05 .00 .00 . 4 .23 .07 .00 .00 10 .2 2 .10 .04 .00 .00 4 , .4 3 .25 .04 .00 .00 9 .8 .13 .08 .87 .00 10 ,.0 .23 .07 .00 .00 10 ,,2

1 .85 .08 .01 .00 2

2 .22 .07 .01 .00 1 . 5

2 .28 .06 .01 . 00 7

2 .39 .08 .02 .00 8

2 .56 .06 .01 .00 7

2 .25 .06 .01 .00 4

2 .23 .08 .01 .00 1 . 6

2 .14 .08 .02 .00 4

2 . 10 .07 .02 .00 2

2 ,, 1 1 .06 .02 .00 2

2 , 08 .08 .01 .00 5

2 ..13 .08 .01 .00 7

2 .,16 .07 .01 . 00 5

2 . 04 .07 .00 .00 5

2 . 02 .07 .00 .00 1 . 0

2 . 05 .07 .00 .00 0

2 . 06 .07 .00 .00 5

2 . 08 .07 .00 .00 5

2 . 17 .07 .03 .00 8

2 . 07 .05 .03 .00 1 . 0

2 . 15 .07 .02 .00 6

2 . 21 .05 .02 .00 7 05 .06 .02 .00 6 . 4 04 .10 .02 .00 1 . 4 04 .09 .02 .00 2 . 0

04 .09 .02 .00 1 . 6

279
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.02 .39 .86 18 .39 10 .16 .04 .10 .02 .00 1 .4 .02 70 .84 18 .96 10 , 1

1

.02 . 1

1

.02 .00 2.5

,02 69 .78 18 .98 9 92 .03 .09 .02 .00 4.8

.02 77 .65 18 .90 9 .89 .03 .08 .01 .00 3 4 .02 ,55 ,80 18 .24 10 .07 .05 .07 .01 .00 3.2

.02 .55 .80 18 .23 10 .05 .05 .08 .02 .00 2.2

.02 .53 .74 18 .12 10 . 17 .05 .08 .02 .00 2.3

.02 .70 .87 19 .55 10 .03 .03 .08 .02 .00 6.6

.02 .72 .79 18 .45 10 .01 .05 .07 .01 .00 4.2

.02 .69 .73 18 .14 10 .15 .05 .06 .01 .00 3.7

.02 ,67 .90 18 .32 9 .94 .05 .07 .01 .00 3.7

.02 .67 .83 18 .57 9 .92 .05 .06 .01 .00 3.9

.03 .89 .88 18 82 10 .35 .02 .07 .01 .00 2.9

.02 .60 .83 18 .51 9 .63 .05 .08 .01 .00 4.8

.02 .50 .81 18 .20 9 .95 .05 . 1

1

.01 .00 .9 .02 .55 .85 18 .15 10 .00 .05 .08 .02 .00 2.6

.02 .52 .84 18 .20 10 . 1

1

.05 .08 .02 .00 2.1

.02 .52 .75 18 .05 9 .98 .05 .07 .01 .00 2.3

.03 1 57 .29 17 .58 9 . 19 .02 .05 .00 .00 2.8

.03 ,95 .22 17 .57 8 .67 .02 .03 .00 .00 6.0

.03 3 23 .39 17 .87 9 .02 .02 .04 .00 .00 4.0

.03 3 .25 .39 17 .58 9 .00 .02 .03 .00 .00 3.2

.04 9 ,55 .34 17 .70 9 .26 .02 .04 .00 .00 3.6

.04 6 54 .37 17 .61 9 .13 .02 .05 .00 . 00 2.8

.03 9 .14 .43 21 .60 1

1

,.67 .02 .20 .00 .00 1 .0 .03 12 00 .49 21 .90 1

1

. 56 .02 .21 .00 .00 1 .8 .03 8 79 .33 18 .61 8 .69 .02 .25 .00 . 00 .4 03 5 , 61 .21 18 .31 8 .65 .02 .25 .00 .00 .4 03 3 .25 .22 18 .04 8 .50 .02 .23 .00 .00 .4 .04 10 .20 .31 20 .49 9 .38 .02 .25 .00 .00 1 .0

. 05 10 .30 .42 21 .01 9 .14 .02 .25 .00 .00 1 .6 .04 1 .45 .36 18 .07 9 . 10 .02 .04 .00 .00 4.0

.03 . 1 .45 .35 18 . 14 9 .88 .02 .04 .00 .00 4.0

.03 2 .04 .59 18 .83 9 .22 .02 .15 .00 .00 1 .0 .04 6 .85 .33 18 .35 9 .19 .02 .03 .00 .00 5.5

.03 7 .07 .40 19 .12 9 . 15 .02 . 15 .00 .00 1 .5 .03 7 .02 .37 20 .08 9 .22 .02 .23 .00 .00 .5 .04 7 .03 .31 14 .74 9 .18 2 .32 .03 .00 .00 2.7

.03 6 .96 .28 14 .82 9 .20 5 .42 .20 .00 .00 .8 .04 6 .91 .34 18 .80 9 .14 .02 .12 .00 .00 2.4

. 04 6 .47 .33 19 .29 9 . 16 .02 .17 .00 .00 1 . 6 .03 6 . 1

1

.36 20 .90 9 .29 .02 .25 .00 .00 1 .6 .04 6 .54 .33 15 .84 9 . 09 1 .91 .05 .00 . 00 2.7

.04 7 .00 .30 15 .78 9 . 18 3 .51 .18 .00 .00 2.0

.03 1 .66 .32 18 .36 8 .94 .02 .10 .00 .00 1 .8 .03 3 .18 .34 18 .99 9 .05 .02 . 15 .00 .00 1 .4 .03 3 .27 .47 17 .79 9 02 .02 .04 .00 .00 3.0

.04 6 .40 .35 18 .51 9 .. 17 .02 . 10 .00 .00 2.6

.04 6 .62 .38 19 .51 9 .06 .02 .15 .00 .00 2.2

.04 6 .31 .34 15 .39 9 ., 14 1 .99 .05 .00 .00 2.8

.04 6 .27 .34 15 . 17 9 ., 14 3 .84 .16 .00 .00 2.6

.04 9 .56 .37 20 .22 9 .07 .02 .27 .00 .00 1 .2 .04 9 .38 .35 18 .98 9 ., 14 .02 .16 .00 .00 1 .7 .04 9 .52 .38 17 .71 9 .. 10 .02 . 1

1

.00 .00 1 .5 .03 3 ,.10 .36 17 .75 9 . 07 .02 .15 .00 .00 .8 .03 3 .28 .47 17 63 9 05 .02 .15 .00 .00 .3 .03 3 . 14 .42 17 .70 9 . 08 .02 .15 .00 .00 .5 .03 1 ,.49 .34 15 .37 9 . 06 2 .03 .03 .00 .00 3.4

.04 3 .05 .35 18 .71 9 ..09 .02 . 1 1 .00 .00 2.

1

.04 6 ,.61 .32 15 .77 9 . 07 3 .98 .18 .00 . 00 2.6

.04 6 .61 .32 17 .77 9 . 10 1 .66 .17 .00 .00 2.5

.03 3 .46 .37 17 .82 9 . 10 .02 .15 .00 .00 .6 .03 3 .20 .34 17 .72 9 . 07 .02 .15 .00 .00 .4

280
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.03 2 .96 .36 18 .13 9 .12 .02 .04 .00 .00 3 . 6 .03 2 .90 .30 18 .79 9 .12 .02 .10 .00 .00 2 .0 .03 3 .16 .42 17 .90 9 .05 .02 .04 .00 .00 3 .0 .03 1 .60 .37 18 .47 9 .03 .02 .12 .00 .00 1 .3 .03 1 .70 .39 19 .55 9 .10 .02 .20 .00 .00 .5 .04 1 .89 .68 23 .21 15 .27 3.48 .10 .00 .00 6 .0 .04 1 .85 .70 23 .07 14 .29 3.97 . 10 .00 .00 6 .8 .04 1 .87 .73 24 .61 14 .34 4.51 .09 .00 .00 5 .6 .04 1 .82 .63 24 .54 13 .85 4.03 .09 .00 .00 3 .6 .03 1 .45 .62 22 .48 15 .78 3.97 .08 .00 .00 2 .4 .02 1 .22 .39 22 .57 14 .97 2.73 .13 .00 .00 5 .2 .03 1 .35 .64 21 .94 16 .38 5.31 .08 .00 .00 .6 .02 1 .20 .94 22 , 10 15 .52 3.38 .12 .00 .00 1 .0 .04 1 .62 .65 24 .37 16 .31 5.02 .10 .00 .00 .0 .03 1 .64 .68 24 .74 14 .02 4.32 .09 .00 .00 7 .4 .04 1 .66 .64 23 .39 14 .46 4.73 .09 .00 .00 2 .0 .03 1 .71 .64 23 .08 14 .41 4.06 .09 .00 .00 6 .0 .04 1 .69 .73 23 .13 13 .73 3.20 . 1

.00 .00 17 .0 .04 1 .68 .68 23 .44 13 .64 3.23 . 10 .00 .00 18 .2 .03 1 .70 .69 23 .09 13 .93 3.10 .10 .00 .00 13 .7 .03 1 .65 .72 24 .73 14 .43 4.53 .09 .00 .00 9 .2 .03 1 .64 .68 24 .74 14 .02 4.32 .09 .00 .00 7 .4 .04 2 .00 .56 19 .85 9 .77 .28 .04 .00 .00 8 .0 .05 1 .91 .51 19 .68 9 .30 .16 .06 .00 .00 7 .8 .05 1 ,.85 .49 19 .95 9 .40 .19 .05 .00 .00 7 .6 .03 2 , 13 .55 24 .30 13 ..41 .22 .09 .00 .00 9 . 4 .02 1 ,.87 .54 19 .98 9 ..70 .16 .06 .00 .00 10 .0 .02 1 .89 .54 20 .03 9 ,.62 .16 .06 .00 .00 9 .4 .02 1 ..87 .55 20 .05 9 . 62 .16 .06 .00 .00 9 .6 .04 1 .99 .57 19 .84 9 .77 .29 .04 .00 .00 8 .7 .03 1 .87 .53 18 .80 12 .,28 2.11 .05 .00 .00 5 .5 .02 1 ..79 .53 20 ..51 9 .,80 .16 .07 .00 .00 10 .0 .03 - 1 .,91 .68 18 .. 1

12 .,46 2.14 .05 .00 .00 3 , .9 .03 1 ..95 .60 19 ,.51 9 .,60 .15 .05 .00 .00 7 .8 .10 1 . 85 .61 24 ..49 12 . 89 .19 . 10 .00 .00 6 .,8

.06 1 .,88 .57 19 ..76 9 . 34 .19 .08 .00 .00 6 .,8 .09 2 . 04 .55 24 ..40 12 . 85 .17 .12 .00 .00 7 . 8 .10 2 . 16 .57 24 ..58 12 . 78 .17 .10 .00 .00 6 . 6 .05 1 .,94 .50 20 ..13 9 . 26 .18 .06 .00 .00 9 ..8 .05 1 . 93 .52 20 . 01 9 . 34 .18 .07 .00 .00 8 . 4

.09 2 . 07 .56 24 .,31 12 . 87 .17 . 1 1 .00 .00 6 . 0 .05 1 . 93 .52 20 .,36 9 . 17 .19 .06 .00 .00 9 . 2 .03 2 . 08 .59 20 . 27 9 . 42 .08 .07 .00 . 00 8 . 5

.03 1 . 95 .53 20 . 12 9 . 18 .30 .08 .00 .00 9 . 2 .03 1 . 85 .52 20 . 04 9 . 27 .30 .08 .00 .00 9 . 2 .03 1 . 84 .53 19 . 99 9 . 33 .30 .08 .00 .00 10 . 6 .10 2 . 1

.63 24 . 53 12 . 83 .17 .12 .00 .00 6 . 8 .03 1 . 79 .50 19 . 90 9 . 47 .30 .10 .00 .00 9 . 8 .07 2 . 04 .53 20 . 04 9 . 35 .30 .07 .00 . 00 7 . 0 .06 2 . 01 .53 19 . 89 9 . 38 .30 .07 .00 . 00 8 . 2 .08 1 . 22 .84 24 . 97 12 . 86 .18 .12 .00 .00 10 . 1 .04 1 . 86 .54 19 . 71 9 . 53 .60 .05 .00 .00 6 . 6 .05 1 . 92 .53 19 . 86 9 . 37 .30 .07 .00 .00 8 . 0

.04 1 . 76 .52 19 . 57 9 . 40 .60 .05 .00 .00 10 . 4 .06 1 . 82 .49 19 . 80 9 . 29 .16 .06 .00 .00 7 . 4 .06 1 . 78 .49 19 . 72 9 . 33 .16 .06 .00 .00 7 . 4 .06 1 . 78 .50 19 . 73 9 . 39 .16 .07 .00 .00 7 . 8

. 10 1 . 92 .54 24 . 58 13 . 12 . 10 .06 .00 . 00 8 . 6 .09 1 . 94 .56 24 . 36 13 . 04 .20 .07 .00 .00 8 . 2 .06 1 . 75 .55 19 . 79 9 . 44 .21 .07 .00 . 00 5 . 6 .06 1 . 73 .56 19 . 71 9 . 32 .20 .07 .00 .00 7 . 3

.09 1 . 82 .51 24 . 19 13 . 08 .20 .07 .00 .00 7 . 0 .05 1 . 83 .54 20 . 35 9 . 29 .16 .06 .00 .00 10 . 9

281
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.06 1 .90 .61 19 .77 9 .29 .16 .05 .00 .00 8 .0 .06 1.92 .54 20 .01 9 .34 .29 .05 .00 .00 7 .2 .06 1 .84 .51 19 .82 9 .29 .16 .08 .00 .00 7 .0 .08 1.83 .72 24 .66 13 .13 .12 .07 .00 .00 9 0 .05 1 .82 .52 20 .41 9 .20 .16 .05 .00 .00 1

1

.6 .08 1 .50 .70 23 .02 12 .65 2 .20 .06 .00 .00 13 .4 .03 2.01 .76 20 .60 9 .50 .14 .08 .88 .00 9 .8 .09 2.22 .78 23 .83 12 .83 .08 . 1

1

.84 .00 9 .2 .09 2.09 .67 24 .58 12 .96 .14 .09 .00 .00 6 .8 .09 1 .93 .52 24 .16 12 .98 .21 .07 .00 .00 6 . 4 .04 1 .23 .84 20 .39 9 .27 .16 .07 .00 .00 12 .5 .05 1 .62 .45 18 .84 13 .05 2 .13 .08 .00 .00 2 . 4 .05 1 .70 .46 18 .73 13 .00 2 .14 .07 .00 .00 1 .5 .03 1 .91 .61 18 .15 13 .54 2 .10 .05 .00 .00 .4 .03 1.82 .50 20 .21 9 .58 .23 .07 .00 .00 10 .2 .03 1 .61 .46 18 .30 12 .55 2 .10 .04 .00 .00 4 . 4

.03 1 .67 .43 18 .97 12 .28 3 .25 .04 .00 .00 9 .8 .03 2.02 .89 20 .04 9 .75 .13 .08 .87 .00 10 .0 .03 1 .82 .50 20 .21 9 .58 .23 .07 .00 .00 10 .2 .03 1 .32 .21 17 .33 12 .37 1 .85 .08 .01 .00 .2 .04 1.54 .22 17 .58 12 .47 2 .22 .07 .01 .00 1 .5 .03 1 .78 .22 17 .73 12 .82 2 .28 .06 .01 .00 .7 .03 2.33 .35 17 .75 12 .78 2 .39 .08 .02 .00 .8 .03 2.43 .22 17 .79 12 .96 2 .56 .06 .01 .00 .7 .03 2.25 .16 17 .65 12 .75 2 .25 .06 .01 .00 .4 .03 1 .43 .26 17 .40 12 .74 2 .23 .08 .01 .00 1 6 .03 1 .56 .42 17 .48 12 .59 2 . 14 .08 .02 .00 . 4 .02 1 .63 .46 17 .46 13 .04 2 .10 .07 .02 .00 .2 .03 1 .56 .43 17 .54 13 .78 2 . 1

1

.06 .02 .00 .2 .03 1 .74 .46 17 .42 12 .20 2 .08 .08 .01 .00 .5 .03 1 .65 .23 17 .34 12 .42 2 . 13 .08 .01 .00 .7 .03 1 .70 .15 17 .30 12 ,.54 2 . 16 .07 .01 .00 .5 .03 * 1.62 .51 17 .48 14 .31 2 . 04 .07 .00 .00 .5 .03 1 .67 .61 17 .34 14 ,.84 2 .02 .07 .00 . 00 1 . 0 .03 1 .60 .35 17 .42 14 .. 45 2 .05 .07 .00 . 00 .0 .03 1 .66 .40 17 . 44 14 .89 2 .06 .07 .00 .00 .5 .03 1 .56 .50 17 .54 14 ,41 2 .08 .07 .00 .00 .5 .03 1 .56 .47 18 .21 13 .80 2 . 17 .07 .03 .00 .8 .03 1 .68 . 44 18 .49 15 .,16 2 ,07 .05 .03 .00 1 .0 .03 1 .75 .43 18 .03 14 . 20 2 .15 .07 .02 .00 .6 .03 1 .63 .44 17 ..97 13 . 68 2 21 .05 .02 .00 .7 .02 .71 .78 18 81 9 . 92 .05 .06 .02 .00 6 . 4 .02 .35 .84 18 ..21 10 . 15 04 . 10 .02 . 00 1 . 4 .02 .37 .84 18 . 57 10 . 16 04 .09 .02 . 00 2 . 0 .02 .37 .84 18 . 37 10 . 1

1

04 .09 .02 .00 1 . 6 .02 .39 .86 18 . 39 10 . 16 04 . 10 .02 .00 1 . 4 .02 .70 .84 18 . 96 10 . 1 1 02 . 1 1 .02 .00 2 . 5

.02 .69 .78 18 . 98 9 . 92 03 .09 .02 .00 4 . 8 .02 .77 .65 18 . 90 9 . 89 03 .08 . 01 . 00 3 . 4 .02 .55 .80 18 . 24 10 . 07 05 .07 .01 . 00 3 . 2

.02 .55 .80 18 . 23 10 . 05 05 .08 .02 .00 2 . 2 .02 .53 .74 18 . 12 10 . 17 05 .08 .02 .00 2 . 3

.02 .70 .87 19 . 55 10 . 03 03 .08 .02 .00 6 . 6 .02 .72 .79 18 . 45 10 . 01 05 .07 .01 .00 4 . 2 .02 .69 .73 18 . 14 10 . 15 05 .06 .01 .00 3 . 7 .02 .67 .90 18 . 32 9 . 94 05 .07 .01 .00 3 . 7 .02 .67 .83 18 . 57 9 . 92 05 .06 .01 . 00 3 . 9 .03 .89 .88 18 . 82 10 . 35 02 .07 .01 .00 2 . 9 .02 .60 .83 18 . 51 9 . 63 05 .08 . 01 .00 4 . 8 .02 .50 .81 18 . 20 9 . 95 05 . 1

1

.01 .00 9 .02 .55 .85 18 . 15 10 . 00 05 .08 .02 .00 2 . 6 .02 .52 .84 18 . 20 10 . 1

1

05 .08 .02 .00 2 . 1 .02 .52 .75 18 . 05 9 . 98 05 .07 .01 .00 2 . 3

282
Appendix I. FN Data Base, continued

c Mn Si Cr Ni N Mo Cb Ti FN

.03 1 .57 .29 17 .58 9 .19 .02 .05 .00 .00 2.8

.03 .95 .22 17 .57 8 .67 .02 .03 .00 .00 6.0

.03 3.23 .39 17 .87 9 .02 .02 .04 .00 .00 4.0

.03 3.25 .39 17 .58 9 .00 .02 .03 .00 .00 3.2

.04 9.55 .34 17 .70 9 .26 .02 .04 .00 .00 3.6

.04 6.54 .37 17 .61 9 .13 .02 .05 .00 .00 2.8

.03 9.14 .43 21 .60 1

.67 .02 .20 .00 .00 1 .0 .03 12.00 .49 21 .90 1

.56 .02 .21 .00 .00 1 .8 .03 8.79 .33 18 .61 8 .69 .02 .25 .00 .00 .4 .03 5.61 .21 18 .31 8 .65 .02 .25 .00 .00 .4 .03 3.25 .22 18 .04 8 .50 .02 .23 .00 .00 .4 .04 10.20 .31 20 .49 9 .38 .02 .25 .00 .00 1 .0 .05 10.30 .42 21 .01 9 .14 .02 .25 .00 .00 1 . 6 .04 1 .45 .36 18 .07 9 .10 .02 .04 .00 .00 4.0

.03 1 .45 .35 18 .14 9 .88 .02 .04 .00 .00 4.0

.03 2.04 .59 18 .83 9 .22 .02 .15 .00 .00 1 .0 .04 6.85 .33 18 .35 9 .19 .02 .03 .00 .00 5.5

.03 7.07 .40 19 .12 9 .15 .02 .15 .00 .00 1 .5 .03 7.02 .37 20 .08 9 .22 .02 .23 .00 .00 .5 .04 7.03 .31 14 .74 9 .18 2 .32 .03 .00 .00 2.7

.03 6.96 .28 14 .82 9 .20 5 .42 .20 .00 .00 .8 .04 6.91 .34 18 .80 9 . 14 .02 .12 .00 .00 2.4

.04 6.47 .33 19 .29 9 ,.16 .02 .17 .00 .00 1 .6 .03 6.11 .36 20 .90 9 ,29 .02 .25 .00 .00 1 .6 .04 6.54 .33 15 .84 9 ,09 1 .91 .05 .00 .00 2.7

.04 7.00 .30 15 .78 9 . 18 3 .51 . 18 .00 .00 2.0

.03 1.66 .32 18 .36 8 ,.94 .02 . 10 .00 .00 1 .8 .03 3.18 .34 18 .99 9 .05 .02 .15 .00 .00 1 .4 .03 3.27 .47 17 .79 9 .02 .02 .04 .00 .00 3.0

.04 6.40 .35 18 .51 9 ..17 .02 . 10 .00 .00 2.6

.04 6.62 .38 19 .51 9 ,06 .02 .15 .00 .00 2.2

.04 6.31 .34 15 .39 9 ..14 1 .99 .05 . 00 .00 2.8

.04 6.27 .34 15 .17 9 .. 14 3 .84 . 16 .00 .00 2.6

.04 9.56 .37 20 .22 9 .,07 .02 .27 .00 .00 1 .2 .04 9.38 .35 18 ,.98 9 ..14 .02 .16 .00 .00 1 .7 .04 9.52 .38 17 ,.71 9 .,10 .02 . 1

.00 .00 1 .5 .03 3.10 .36 17 .75 9 . 07 .02 .15 .00 .00 .8 .03 3.28 .47 17 ,.63 9 . 05 .02 .15 .00 .00 .3 .03 3.14 .42 17 ,.70 9 . 08 .02 .15 .00 .00 .5 .03 1 .49 .34 15 ,.37 9 . 06 2 .03 .03 .00 . 00 3.4

.04 3.05 .35 18 ,.71 9 . 09 .02 . 1

.00 .00 2.

.04 6.61 .32 15 .,77 9 . 07 3 .98 .18 .00 .00 2.6

.04 6.61 .32 17 .,77 9 . 10 1 ,.66 .17 .00 .00 2.5

.03 3.46 .37 17 ,,82 9 . 10 .02 .15 .00 .00 . 6 .03 3.20 .34 17 ,.72 9 . 07 ,02 .15 .00 .00 . 4 .03 2.96 .36 18 , 13 9 . 12 ,02 . 04 .00 . 00 3.6

.03 2.90 .30 18 . 79 9 . 12 .02 . 10 .00 .00 2.0

.03 3.16 .42 17 . 90 9 . 05 02 . 04 .00 . 00 3.0

.03 1.60 .37 18 . 47 9 . 03 02 .12 .00 . 00 1 .3 .03 1 .70 .39 19 . 55 9 . 10 02 .20 .00 .00 .5

283



INFLUENCE OF MOLYBDENUM ON THE STRENGTH AND

TOUGHNESS OF STAINLESS STEEL WELDS FOR CRYOGENIC SERVICE*

C. N. McCowan and T. A. Siewert

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

E . Kivineva

Department of Metallurgy

University of Oulu

Oulu, Finland

Molybdenum additions to austenitic stainless welds were found

to increase the 4-K yield strength by approximately 30 MPa/wt.%.

Molybdenum additions had little effect on the 76 -K Charpy V-Notch

impact energy, with one exception: when the molybdenum content

was raised from 1.7 to 3.8 wt.% in an otherwise equivalent stainless steel composition of approximately 17Cr—9Ni—6 . 6Mn—0 . 17N , the

absorbed energy decreased from 33 to 16 J. At a nickel content of

14 wt.%, the higher molybdenum contents did not reduce the impact

toughness. The loss in impact toughness for the 9 -wt.% nickel,

3. 8 -wt.% molybdenum composition was linked to an increase in both

small and large inclusions in the weld. The effects of nickel

and manganese on the cryogenic strength and toughness are also

reported.

INTRODUCTION

Recently there has been an interest in the development of higher

strength austenitic stainless steels for lower cost and more compact

magnetic fusion structural designs. Fusion magnet components are expected

to operate at temperatures near 4 K. Austenitic stainless steels are one of the common structural materials chosen for service in this environment

because, unlike many high strength materials, these alloys have good

toughness at temperatures approaching absolute zero. For example, high

*Accepted for publication in Welding Metallurgy of Structural Steels , Metallurgical Society of AIME, Warrendale, Pennsylvania.

285
strength type 316LN stainless steel base materials have 4-K yield strengths

of 1020 MPa, 76 -K Charpy V-notch absorbed energies of 89 J, and 4-K fracture

toughness [Kj c (J)] values of 224 MPaVm. 1 The properties of these nitrogenstrengthened base materials cannot be totally utilized, however, until welds

of comparable toughness at these elevated strengths are also developed.

Nitrogen additions can increase the yield strength of 300- series

stainless steel weld metals to levels comparable to those of the nitrogenstrengthened base materials. The problem is that toughness values for the

welds at cryogenic temperatures have been consistently lower than those for

base materials of equal strength. 2 For this reason, other strengthening

elements that have a less detrimental effect than nitrogen on toughness are being investigated.

The effect of molybdenum on the strength and toughness in shielded

metal arc (SMA) welds is reported in this study, and interactive effects

with nickel, nitrogen, and manganese are also investigated.

BACKGROUND

Much of the data pertaining to the effects of alloying element additions on the strength and toughness of austenitic stainless steels at 4 K

are from studies of base materials. In types 304 .and 316 stainless steel

base materials, the strength—toughness relationship remains essentially the

same, whether the strength is increased by nitrogen, carbon, or molybdenum

additions. 3 ’ 4 The toughness decreases as the strength increases. From this

perspective, nitrogen, which increases the 4-K yield strength in both 300-

series weld and base metal by approximately 3000 MPa/wt.%, would therefore

be expected to have a more detrimental effect on toughness than molybdenum. 5

' 8 Molybdenum reportedly increases the yield strength at 4 K by

only 40 to 80 MPa/wt.% in 300 series base materials. 3 ’ 6 ’ 9 The only alloying element that has been found to improve the strength-toughness relationship in 300 series stainless steels at cryogenic temperatures is nickel.

Increased nickel contents increase the toughness while having very little

effect on the strength. 10

Although, substitutional alloying by molybdenum in 300-series stainless steel base materials has been found to increase strength and decrease

toughness, as molybdenum contents are increased, unexpected decreases in

toughness have been observed. 3 One cause for the loss of toughness in

molybdenum- alloyed stainless steels is the formation of molybdenum- rich

chi and laves phases. Formation of these brittle intermetallic phases

reduced the 76-K Charpy V-notch energy by more than 100 J in a 12Cr-12NiO.IN-Mo base metal alloy when the molybdenum content was increased from 5 to 7 wt.%. 9 However, not all reported toughness losses for molybdenumalloyed stainless steels can be clearly related to embrittlement by these

intermetallic phases. A type 316 stainless steel base metal, for example,

had unexpectedly low 4-K fracture toughness values at 3.8-wt.% molybdenum

contents. 3 Presumably, the molybdenum had not been uniformly distributed

in this alloy. A second phase was present in the microstructure, but it was not rich in molybdenum.

286
In welds, nonuniform distributions of alloying elements are common

owing to the dendritic solidification process. Molybdenum, in particular,

segregates extensively in austenitic stainless steel welds. 11 Therefore,

decreases in weld metal toughness, similar to the decreases observed in base

metals, may also result from an inhomogeneous distribution of molybdenum.

In the case of a 25Cr-16Ni-0 . 34N-Mo weld, ductilities were found to decrease

rapidly after the molybdenum content exceeded 3.5 wt.%. 12

MATERIALS AND PROCEDURE

The test matrix was designed to evaluate welds with molybdenum contents

of approximately 0.02, 2, and 3.8 wt.%. At each level of molybdenum content, there were two possible nickel and nitrogen levels. In addition,

several manganese additions were included in the test matrix. The alloying

elements were added to the shielded metal arc (SMA) electrode coatings. The

weld compositions were determined by optical emission spectroscopy, and the

ferrite number (FN) of the weld deposits was measured magnetically; both are given in table 1.

Test welds were produced by using a series of 3 . 2 -mm-diameter electrodes that had core wires from a single heat of type 308 stainless steel.

The base material was 13-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 welding,

as specified in AWS A5.4-78. The same welder, power supply, and welding

parameters (110 A, 22 V, identical bead sequences, and a heat input of

approximately 1.6 kJ/mm) were used to minimize variations in the weldments.

The interpass temperature was maintained at 93°C.

Two all-weld-metal 6 -mm-diameter 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 (in liquid helium) with a strain rate of

9.5 x 10" 5 s' 1

. A 25-mm gauge length was used. Specific details on the 4-K

tensile testing equipment and procedures have been reported previously. 13

Five 10-mm-square Charpy V-notch 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 (in liquid nitrogen)

Scanning electron microscopy was employed in the acquisition of inclusion distribution data from polished weld metal samples. In addition,

wavelength dispersive spectroscopy (WDS) was used to evaluate molybdenum

segregation between the ferrite and austenite phases, and energy dispersivespectroscopy (EDS) was used to evaluate the composition of the inclusions on the fracture surfaces.

Welds 6 and 13 were evaluated with x-ray diffraction techniques. They

were step-scanned at 0.02-degree intervals of 29 for 25 s/step. The tube

voltage was 44 kV at 40 mA. The radiation was CuK/? (ACu/3 - 1.39225 A).

287
RESULTS AND DISCUSSION

Strength

The yield strength of the weld metal at 4 K increased as a function of

molybdenum, manganese, and nitrogen content. Relationships between alloy

content and the ultimate strength could not be resolved. The strengths

determined for the welds are listed in table 2.

The ultimate strength of the welds had a mean value of 1400 MPa, which

is consistent with previous results obtained for SMA weld metals with compositions similar to the 9-wt.% nickel welds tested in this study. 5 Inhomogeneities in the weld structures and variations in ferrite contents evidently

contributed to the scatter in the ultimate strength values of this study.

To evaluate the contributions of the alloying elements to the yield

strengths of these welds, the data were analyzed with stepwise linear

regression techniques. The 4-K yield strength (in MPa) was best predicted

by the equation:

Yield Strength = 180 + 3200[N] + 33[Mo] + 32[Mn] + 13[Ni] (1)

This equation (elements in wt.%) had an F value of 920, a coefficient of

regression (R2

) of 0.98, and a standard error of estimate of 31 MPa.

As expected, the interstitial strengthening of nitrogen had a much

greater effect on the yield strength at 4 K than the solid- solution

strengthening contributions of molybdenum, manganese, or nickel. Equation 1 shows the strengthening of nitrogen (3200 MPa/wt.% N) to be two orders of

magnitude greater than that of the solid- solution elements. This nitrogen

coefficient agrees well with those previously determined for both weld and

base metals. 6-8 The effect of nitrogen on the yield strength at 4 K is

shown in figure 1.

Much of the variation in yield strength, however, can be attributed to

the solid-solution strengthening of molybdenum and manganese. To show the

strengthening contributions made by these elements more clearly, the yield

strength data were normalized to a 0.05-wt.% nitrogen content using the

3200 MPa/wt.% strengthening coefficient from equation 1 (figure 2). The

slopes of the dashed lines in figure 2 indicate the increases in yield

strength due to molybdenum additions at the three manganese contents present

in this study. The difference in the yield strength between respective

lines indicates the strengthening by manganese. Clearly, molybdenum and

manganese also contribute to the 4-K yield strength of these welds.

Although both figure 2 and equation 1 indicate the relative magnitude

of the strengthening contributions made by the alloying elements, singular

strengthening coefficients to represent the effect of molybdenum or manganese could not be defined. The coefficients for these elements did not

remain constant over the compositional range considered in this study.

288
For example, when the 14-wt.% nickel data were evaluated separately, the

molybdenum coefficient decreased to 25 and the manganese coefficient

decreased to 15. Further manipulation of the data indicated that the

substitutional strengthening effects of manganese and molybdenum may

decrease as the nickel content, nitrogen content, or both increase. The

coefficients in equation 1, therefore, represent the average effects of

these elements.

Toughness

The Charpy V-notch absorbed energy (AE) and lateral expansion (LE)

values determined for the welds are reported in table 2. Impact toughness

at 76 K decreased as the interstitial alloy content of the welds increased.

The effects of substitutional alloy content on the impact toughness varied.

Nickel increased both AE and LE for alloys having low nitrogen contents

(figures 3 and 4). At higher nitrogen levels (> 0.15-wt.% N) , however, the

beneficial effect of increased nickel content on the impact toughness was

reduced.

The effect of molybdenum on the impact energy was related to the nickel

content. At nickel contents of 14 wt.%, molybdenum did not adversely affect

either the AE or LE. Increased molybdenum contents, in fact, may have been

slightly beneficial to the toughness in the higher nickel content welds. In

the 9 -wt.% nickel welds, however, the AE was much lower for the 3. 8 -wt.%

molybdenum alloy (weld 6). To verify this finding, a new weld was made with

another electrode of similar composition. Again, low AE values (17 J) were

found for the 4Mo—9Ni type stainless steel weld metal. These AE values are approximately ' 50 percent lower than that of the 2Mo—9Ni alloy (weld 13).

Although the chromium content in welds 6 and 13 also differs, otherwise

they have quite similar compositions and the decrease in AE is felt to be

directly related to 3. 8 -wt.% molybdenum content. Lower molybdenum contents

had no adverse effect on impact toughness of the 9-wt.% nickel welds.

Possible explanations for the decrease in toughness associated with the

3. 8 -wt.% molybdenum welds are given later in this report.

Regression analysis was applied to the AE and LE data, but equations

that were significant at the 99% confidence level could not be developed.

Nevertheless, the analysis consistently suggested that: (1) nitrogen re- duced the AE and LE, (2) nickel increased the AE and LE, and (3) manganese

and ferrite decreased the LE. Because the chromium content was lowered as the molybdenum content was increased to maintain a constant ferrite potential, the effects of these elements on toughness is uncertain. With the

exception of the 4Mo—9Ni weld, however, the effect of molybdenum on toughness was small in comparison with the nitrogen effect. The decrease in LE

associated with manganese indicated here agrees well with a previous studv

performed on 9-wt.% nickel SMA welds. 6 In general agreement with past

studies, nitrogen3 and ferrite 14 reduce the toughness, and nickel 10 increases the toughness

289
Evaluation of the 3.8-wt.% Molybdenum Alloy

In an effort to determine the reason for the low impact toughness of

the 4Mo—9Ni alloy (weld 6) , the microstructure of the weld and the fracture

surfaces of the test specimens were evaluated and compared with those of

weld 13. As previously mentioned, welds 6 and 13 had very similar compositions with the exception of their chromium and molybdenum contents. The

reduction in absorbed energy from 33 J (weld 13) to 16 J is attributed to

the increase in molybdenum content, from 1.7 wt.% in weld 13 to 3.8 wt.% in

weld 6.

The two welds were found to have similar ferrite contents and primary

ferritic solidification morphologies during light microscope evaluations of

mechanically polished and etched weld metal specimens. The extent of molybdenum segregation between the ferritic and austenitic phases of welds 6 and

13 was also similar. Line scans using wavelength dispersive spectroscopy

showed molybdenum contents to be approximately 1 wt.% higher in the ferritic

phase than in the austenitic matrix.

Evaluations of the weld metals to determine the presence of intermetallic phases were inconclusive. No evidence of intermetallic phases was found during light microscope examination of the respective weld metals.

X-ray analysis showed that both welds 6 and 13 had very similar types and

quantifies of phases present in their structures. These findings would

indicate that the lower absorbed energy found for the 3. 8 -wt.% molybdenum

alloy cannot be attributed to the formation of brittle intermetallic phases.

The only resolvable differences found between the two welds were

related to inclusion sizes. Large inclusions were found on the Charpy Vnotch fracture surfaces of the 3.8-wt.% molybdenum weld. These inclusions

(figure 5) had (a) diameters ranging from 20 to 100 pm; (b) compositions

rich in calcium, manganese, silicon, and titanium; and (c) glassy, conchoidal fractures associated with them. Although large inclusions were also

found on lower molybdenum content fracture surfaces, they were observed less

frequently and generally were not as large. In addition to finding large

inclusions on the fracture surfaces of weld 6 specimens, inclusion counts on polished specimens showed a difference between the size distributions for

the smaller inclusions in weld 6, figure 6. Weld 6 was had a greater number

of small diameter inclusions than the other welds. Whether the presence of

more small and large inclusions in alloy 6 was solely responsible for lowering the toughness in these samples is not known.

Inclusions have, however, long been recognized as void nucleation sites

in materials having dimpled rupture fracture modes , and they do affect the

resistance of engineering alloys to fracture. 15 Attempts to correlate inclusion content and fracture toughness [Kj c (J)] at cryogenic temperatures

are not directly applicable to the Charpy V-notch toughness results of this

study, but they indicate that as inclusion spacing decreases, the toughness

decreases. 16 Although more complete investigations of inclusion morphology

and spacing were beyond the scope of this study, ductile dimple density

counts taken on the fractured tensile specimen surfaces of welds 6 and 1 support the argument that decreased inclusion spacing lowers the toughness.

290
The weld 6 specimens had a substantially higher dimple density (4.4 x 10 s dimples/mm2

) than the weld 1 specimens (3.2 x 10 5 dimples/mm2 ). This

indicates that the inclusions nucleating the voids in weld 6 were more

closely spaced. Further study is needed to determine the quantitative relation between the presence of the large number of small inclusions (figure 6)

observed in weld 6 and inclusion spacing. The very large inclusions, found

on weld 6 fracture specimen surfaces, could also be speculated to reduce

toughness. As a general observation, when the inclusion size is increased,

the strain required to nucleate a void is decreased. 17 ’ 18

CONCLUSIONS

1. The solid-solution strengthening contribution of molybdenum to the welds

at 4 K is approximately 30 MPa/wt.%.

2. The effect of molybdenum on the 76-K Charpy V-notch absorbed energy and

lateral expansion is small at molybdenum contents of 2 wt.% or less.

3. In 18Cr—9Ni stainless steel SMA welds, molybdenum contents up to 3.8 wt.%

cause substantial decreases in Charpy V-notch absorbed energy and lateral

expansion at 76 K.

4. In 18Cr—14Ni stainless steel SMA welds, increasing the molybdenum content

from approximately 0 to 3.8 wt.% did not adversely affect the Charpy Vnotch toughness at 76 K.

ACKNOWLEDGMENT

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

Fusion Energy.

REFERENCES

1. Mazandarany, F. N. , Parker, D. M. , Koenig, R. F. , Read, D. T.,

A Nitrogen- Strengthened Austenitic Stainless Steel for Cryogenic Magnet

Structures, "Advances in Cryogenic Engineering—Materials," vol. 26,

Plenum Press, New York, 58—170 (1980).

2. McHenry, H. I., Read, D. T., Steinmeyer, P. A., Evaluation of Stainless

Steel Weld Metals at Cryogenic Temperatures, "Material Studies for

Magnetic Fusion Energy Applications at Low Temperatures—II," NBSIR 79-

1609, National Bureau of Standards, Boulder, Colorado, 299 (1979).

3. Purtscher, P. T. , Reed, R. P. , Effect of Molybdenum on the Strength

and Toughness of Stainless Steels at Cryogenic Temperatures, to be

published in "Advances in Cryogenic Engineering—Materials," vol 34,

Plenum Press, New York.

291
4 . Tobler, R. L. , Read, D. T. , Reed, R, P., Strength/Toughness Relationship for Interstitially Strengthened AISI 304 Stainless Steels at 4 K,

"Fracture Mechanics: Thirteenth Conference," Richard Roberts, Ed.,

ASTM STP 743, American Society for Testing and Materials, Philadelphia,

250-268 (1981).

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

Nitrogen in Stainless Steel SMA Welds for Cryogenic Service, Welding

Journal, in press (1987).

6. Yamamoto, S., Yamagami , N. , Ouchi , C., Effect of Metallurgical Variables on Strength and Toughness of Mn—Cr and Ni—Cr Stainless Steels at

4.2 K, "Advances in Cryogenic Engineering—Materials, vol. 32, Plenum

Press, New York, 57-64 (1986).

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

Stainless Steels with Nitrogen and Carbon, "Advances in Cryogenic Engineering—Materials, vol. 30, Plenum Press, New York, 127—136 (1984).

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

at 4 K, "Material Studies for Magnetic Fusion Energy Applications at

Low Temperatures—IX," NBSIR 86-3050, National Bureau of Standards,

Boulder, Colorado, 27-42 (1986).

9. Miura , R. , Development of High Strength Structural Alloys for Superconducting Magnets in Fusion Reactor, unpublished.

10. Reed, R. P., Purtscher, P. T. , and Yushchenko, K. A., Nickel and Nitrogen Alloying Effects on the Strength and Toughness of Austenitic Stainless Steels at 4 K, " Advances in Cryogenic Engineering—Materials,"

vol. 32, Plenum Press, New York, 43—50 (1986).

11. Ogawa, T., Aoki , S., Sakamoto, T. , and Zaizen, T., The Weldability of

Nitrogen-Containing Austenitic Stainless Steel: Part I —Chloride

Pitting Corrosion Resistance, Welding Journal 62(5), 139s-148s (1982).

12. Sakamoto, T. , Abo, H. , Okazaki, T. , Ogawa, T. , Ogawa, H., Zaizen, T.

,

Corrosion Resistant Nitrogen-Containing Stainless Steels for Use by

the Chemical Industry, "Alloys for the Eighties," Climax Molybdenum

Company, Leadville, Colorado, 269-279 (1980).

13. Read, D. T. , Tobler, R. L. , Mechanical Property Measurements at Low

Temperatures, "Advances in Cryogenic Engineering—Materials," vol. 28,

Plenum Press, New York, 17—28 (1982).

14. Szumachowski , E. R. , Reid, H. F. , Cryogenic Toughness of SMA Austenitic

Stainless Steel Weld Metals: Part 1 —Role of Ferrite, Welding Journal

57(11), 325s—333s (1978).

15. Van Stone, R. H. , Cox, J. R. , Low, J. R. , Jr., Psioda, J. A., Microstructural Aspects of Fracture by Dimpled Rupture, International Metals

Reviews 30(4), 157-179 (1985).

292
16. Reed, R. P., Simon, N. J., Strength and Toughness of AISI 304 and 316

at 4 K, Journal of Nuclear Materials 44, 141—143 (1986).

17. Tauoka, J. P., Pampillo, C. A., Low, J. R. , Jr., in "Review of Develop

ments in Plane Strain Fracture Toughness Testing," W. F. Brown Jr.,

Ed., STP 463, American Society for Testing and Materials, Philadelphia

191 (1970).

18. Van Stone, R. H. , Merchant, R. H. , Low, J. R.

Fracture Toughness—Cryogenic Behavior," C. F.

Broadwell, Eds., STP 556, American Society for

Philadelphia, 93 (1973).

Jr.in "Fatigue and

Hickey and R. G.

Testing and Materials,

293
Weld

12

5

6

12

13

28

29

60

23

24

25

26

27

68

69

70

71

Table 1. Weld Compositions and Ferrite Contents.

c Mn Si P S Cr Ni Mo N

FN

Plate

0.033 1.57 0.29 0.013 0.006 17.58 9.19 0.02 0.047 4.9

0.034 1.49 0.34 0.013 0.006 15.37 9.06 2.03 0.034 4.4

0.039 6.31 0.34 0.013 0.007 15.39 9.14 1.99 0.047 3.6

0.036 6.27 0.34 0.013 0.007 15.17 9.14 3.84 0.162 2.9

0.037 6.62 0.38 0.013 0.008 19.51 9.06 0.02 0.151 3.6

0.035 6.61 0.32 0.013 0.007 17.77 9.10 1.66 0.166 3.5

0.028 3.23 0.33 0.014 0.007 19.56 *9.31 0.02 0.186 1.4

0.029 3.21 0.34 0.019 0.007 18.18 9.24 2.17 0.179 1.8

0.036 2.38 0.29 0.024 0.005 19.89 9.40 0.02 0.187 0.9

0.037 3.10 0.29 0.012 0.011 19.51 14.61 0.03 0.044 0.4

0.038 3.01 0.30 0.016 0.010 14.29 14.29 2.01 0.039 0.5

0.040 2.97 0.28 0.020 0.010 15.68 14.33 3.90 0.036 0.4

0.040 6.51 0.29 0.017 0.013 17.80 14.49 1.92 0.038 0.0

0.044 6.71 0.35 0.021 0.013 15.94 14.57 3.84 0.042 0.4

0.028 3.20 0.27 0.012 0.010 19.36 14.67 0.02 0.191 0.3

0.031 3.21 0.27 0.013 0.009 18.76 14.69 1.06 0.182 0.3

0.028 3.10 0.27 0.015 0.009 18.03 14.70 2.02 0.183 0.3

0.027 3.05 0.25 0.018 0.009 16.57 14.74 3.98 0.178 0.4

Table 2. Mechanical Properties of the Welds.

Yield Ultimate Absorbed Lateral

Alloy Strength Strength Energy Expansion

MPa (4 K) MPa (4 K) J (76 K) mm (76 K)

1 460 1598 46 0.69

2 523 1570 39 0.71

5 665 1489 37 0.53

6 1169 1432 16 0.13

12 1027 1376 37 0.31

13 1129 1409 33 0.23

28 999 1335 41 0.38

29 1076 1520 41 0.41

60 998 1227 37 0.41

23 638 1152 51 0.84

24 675 1269 60 0.89

25 752 1359 43 0.84

26 758 1299 59 0.91

27 836 1385 56 0.81

68 1093 1413 34 0.36

69 1098 1493 33 0.38

70 1086 1471 37 0.38

71 1124 1453 42 0.43

294
<0

Figure 1

Nitrogen, wt. %

The effect of nitrogen on the 4-K yield strength.

ct

Figure 2. The 4-K yield strength data, normalized (3200 MPa/wt.%

nitrogen) to a 0.05-wt.% nitrogen content to show the

strengthening contributions of molybdenum and manganese.

295

18

ksi
76-K

Lateral

Expansion,

mm

Figure 3. The effect of molybdenum and nickel content

on the 76-K Charpy V-notch absorbed energy.

0.04

0.03

0.02

0.01

Figure 4. The effect of molybdenum and nickel content

on the 76-K Charpy V-notch lateral expansion.

296

Inches
Number

of

Inclusions

per

mm

Figure 5. Fractograph (SEM) of a Charpy V-notch specimen from weld 6 showing one of the large inclusions on the fracture surface.

Figure 6. Inclusion size distribution results of

welds 1, 6, 12, and 27 acquired at 3000x.

297



HIGH- ENERGY- BEAM WELDING OF TYPE 316LN STAINLESS

STEEL FOR CRYOGENIC APPLICATIONS*

T. A. Siewert, D. Gorni , t and G. Kohn+ Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

Laser and electron-beam welds in 25-mm-thick type 316LN

stainless steel containing 0.16 wt.% N were evaluated. Their

mechanical properties are compared with those of welds produced

with more conventional, lower productivity processes, such as shielded-metal-arc and gas-metal-arc welding. Because these

welds are so narrow (about 2 mm) , tensile tests were performed

on transverse tensile specimens at 4 K. For both welding processes, the fractures occurred in the base metal at a strength

level near 950 MPa, indicating that the weld and heat-affected

zone has a strength similar to that of the base metal. The 4-K

weld fracture toughness averaged near 150 MPa^/m for both welding

processes, a value only slightly less than that for the base

metal and comparable to the best values achieved with conventional welding processes using type 316LN weld metal of similar

strength. The Charpy V-notch absorbed energies averaged near

70 J at 76 K. Metallographic analysis revealed cellular, fully

austenitic solidification with little porosity and no evidence

of hot cracking.

INTRODUCTION

Fusion energy applications demand superior properties from structural

materials. For magnet casings, the material must be stable at 4 K and

possess high strength and toughness. These requirements are met by type

316LN stainless steel. 1 Unfortunately, matching composition welds produced

by conventional processes (shielded metal arc, gas metal arc, and flux conarc welding) generally have a substantially poorer strength and toughness

Submitted to Advances in Cryogenic Engineering—Materials , vol. 34.

truest scientist from Armament Development Authority (RAFAEL), Haifa, T r; r: • +Guest scientist from Israel Atomic Energy Research Center, Beersheva,

Israel

299
combination. 2 To improve the relationship between strength and toughness,

the factors influencing this relationship are now being investigated.

The strength and toughness combination can be improved by modifying the

composition, but this approach introduces the potential for other material

problems, such as (1) a change from primarily ferritic to primarily aus- tenitic solidification, which is more sensitive to hot cracking and (2) the

introduction of strain due to the differing thermal expansion of the weld

and the plate material. One way to change the strength and toughness

combination without changing the material is to choose a different joining

process, especially one that produces a lower inclusion content. 3

Both the electron beam (EB) and laser welding processes are capable of

producing welds with a lower inclusion content. They remelt the base metal,

which initially has a lower inclusion content than more conventional welds.

Also, EB welds are produced in a vacuum and laser welds can be produced

in an inert gas. These conditions nearly eliminate the inclusion-forming

elements that may exist within, or may be aspirated into, the arc atmosphere

of conventional welding processes.

Both the EB and laser welding processes have the advantages of highenergy -density beams. The higher travel speeds and deeper penetration of

the high-energy-density beams results in higher productivity. Associated

with the rapid travel is a narrower weld and heat-affected zone (HAZ) , which

minimizes the effect of welding on the material properties.

The energy density is also a disadvantage: the molten weld metal

reaches a very high temperature at which the alloying elements with high

vapor pressures will evaporate. In type 316LN stainless steel, the elements

most likely to be lost are nitrogen and manganese. Since solid- solution

strengthening by nitrogen has a major role in determining the strength, a significant loss in nitrogen content would result in the loss of the loadbearing capability of the joint. 4 ’ 5 A significant loss in manganese by

evaporation could result in porosity or reduced strength in the weld, since

the manganese raises the solubility of the nitrogen in the steel, and to a lesser degree, substitutionally strengthens the alloy. 4 Therefore, both the

strength and porosity were carefully evaluated in this study.

Besides interacting with the beam to modify the mechanical properties,

the nitrogen content can also affect the weldability. Brooks reported that

the 0.3 wt . % nitrogen in an Fe—21Cr—6Ni—9Mn is directly responsible for

undesirable weld features, such as spatter and porosity, and that the

occurrence of these undesirable features is proportional to the nitrogen

content. 6 Thus, the porosity may not be due only to a loss of manganese,

which reduces the nitrogen solubility, but also to the effect of the highenergy beam on the dissolved nitrogen.

Electron-beam and laser welding differ in their degree of development,

total beam power, and energetic particles (electrons versus photons).

300
Electron-beam welding was developed first and is capable of welding thicker

material. It has met the stringent requirements of pressure vessel fabrication in two-pass welding of 200-mm- thick steel. 7 Also, it has been applied

to the welding of stainless steels of 80-mm thickness in an application for

high- temperature service. 8

Laser welding has not been developed to the power density of some EB

systems, but it is capable of joining at least 18-mm- thick steel in a single

pass. 9 Unlike EB welding, laser welding does not require a vacuum chamber

or even a local vacuum. Nyilas found that laser welding was the most costefficient process for joining 1-mm-thick stainless steel for a superconducting cable conduit. 10

MATERIAL

The heat used in this study and other heats of AISI type 316LN base

metal have been characterized previously in a study that developed predictive equations for the strength and toughness of type 316LN alloys. 5 Data

for this heat agreed closely with the measured values for the other heats,

confirming that this heat is representative of type 316LN base metal. Data

from these tests are included in table 1 along with data on welds produced

with conventional processes. Because the strength and toughness are interrelated, as shown by the trend lines in figure 1, the table also includes a quality index, QI , which normalizes the effects of varying strength and

nickel content and enables easy comparison of the base metal and welds at

various strength levels.

Microstructural analysis confirmed that the steel had a fully austenitic structure. Its composition is given in table 2.

PROCEDURE

Rectangular pieces machined from the plate and cleaned with acetone

were oriented in the butt-joint configuration shown in figure 2. Two welds

were produced by using each of the two welding processes and the parameters

listed in table 3 to ensure an adequate amount of weld metal for the evaluation. Because the laser had insufficient power for full penetration of the

25-mm thickness, a two-pass procedure (one from each side) was used, with

both laser passes produced in the same direction. During the EB procedure

development, frequent gas eruptions occurred in the weld pool. For this

reason, the welding speed was decreased so the gases could escape.

After radiography to evaluate weld quality, standard mechanical property specimens with a weld cross section were removed from the two types of

weldments; specimen orientations are shown in figure 3. The specimens con- sisted of

• two compact (CT) specimens (CT-1 and CT-2) notched with and

against the direction of welding

301
• four Charpy V- notch (CVN) specimens (C-l through C-4) notched

with and against the direction of welding and located near

the top and bottom of the weld

• two transverse tensile specimens (T-l and T-2) located near

the top and bottom of the weld

• a metallographic specimen (M-l) for weld dimension, microstructure, and inclusion measurements.

The CT specimens had a thickness of 25 mm after removing the weld reinforcements and were produced according to ASTM specification E 399 with the

notches carefully centered in the weld metal. The compact tension specimens

were tested at 4 K by using the single-specimen unloading compliance technique described in ASTM E 813 to determine the Kj c (J) fracture toughness.

The notch region was precompressed one percent of total thickness to promote

a straight fatigue crack. Both sides were grooved after fatigue precracking

to avoid possible propagation of the fatigue crack into the HAZ or base

metal. The techniques used to measure the fracture toughness at 4 K have

been described elsewhere. 12

Charpy V-notch specimens conforming to ASTM specification E 23 were

produced with their notches centered in the weld. The directional nature of

welding results in a directional microstructure. The effect of this directionality on the toughness was measured by notching the CT and CVN specimens

so the cracks would propagate with the direction of welding in some of the

specimens and against the direction in others.

The tensile specimens conformed to ASTM specification A 370. Contrary

to the usual procedure for welds, they had a transverse orientation. With

conventional welding processes, an all-weld-metal specimen, 6 mm in diameter, can be removed along the length of the weld. With the EB process, 2-mm

width of the welds precludes measurement of only the weld metal, and so a transverse geometry was chosen to measure the joint efficiency. The cross

head speed was 0.2 mm-min 1

. The techniques used to measure the strength at

4 K have been described elsewhere. 13

The metallographic section was examined before and after etching to

enable the mechanical properties to be related to the microstructure and

inclusion density. The fracture surfaces of the CT specimens were also

examined for inclusion density. These weldments and undeformed specimens

were measured magnetically according to the procedure of AWS A4.2 for

determination of the delta- ferrite content.

RESULTS AND DISCUSSION

Tensile Tests

The transverse tensile data for the two welds are shown in table 4.

The yield and tensile strengths are close to those of the base metal (listed

in table 1). This similarity in strength was obvious when the specimens

were examined: each broke outside the weld in the base metal, which is an

302
indication of a weld strength greater than the base metal strength. The EB

weld specimens broke about 3 mm from the weld fusion line, and the laser

weld specimens broke near the specimen shoulder. Since the weldments were

limited in strength by the base metal, nitrogen and manganese loss from the

molten weld puddle due to vaporization had a negligible effect on the joint

strength. The measured nitrogen content of the EB weld was 0.135 wt . % , a decrease of 0.015 wt.% from the base metal analysis listed in table 2.

The fracture of the EB weld specimens so near the fusion line indicates

a weak HAZ. However, examination of the fracture surface revealed the

fracture initiated at a large inclusion, which could have determined the

fracture location. Despite these weaknesses, the weld strength still fell

within the range measured for type 316LN base metal in table 1. Its elongation was poorer than expected because this was the short transverse direction in the base metal plate.

Charpy V-notch Tests

The resistance to fracture of CVN specimens at 76 K is given in

table 5. The difference between specimens with cracks running with and

against the welding direction was insignificant. From the convergence of

grain boundaries and possible brittle interdendritic regions in the direction of weld solidification, we would expect cracks running in the direction

of welding to have slightly lower absorbed energy and thus to be the more

conservative measurement of weld toughness.

The laser welds had a 19 percent lower absorbed energy than the EB

welds. This difference is attributed to differences in welding parameters

rather than to the type of process. However, both welds had twice the

absorbed energy of conventional welds; see the representative data in

table 6. The significantly higher absorbed energy indicates a substantial

improvement in fracture resistance.

Compact Tension Tests

The fracture toughness values at 4 K included in table 7 are substantially larger than those for the conventional processes in table 8 and are within 10 percent of the data for the base metal in table 1. A better

comparison can be made by evaluating the strength—toughness combination in

the form of a Quality Index (QI) , the product of strength and toughness.

The QI values for the laser and EB welds are within 10 percent of that for

the base material (table 1), whereas the QI values for conventional welds

average 30 percent lower (table 8).

Microstructure

The radiographic and metallographic evaluation of the laser weld

revealed significant porosity (up to 1 mm in diameter), which is unacceptable in a production weld. No porosity was seen in radiographic evaluation

of the EB weld, but some microporosity was seen in the metallographic

evaluation. The difference in porosity is attributed to differences in the

welding parameters, not the type of process. If the laser weld process wer303
repeated using the lower travel speed of the EB weld prooess, porosity would

not be expected. The porosity probably contributed to the lower toughness

of the laser weld.

The metallographic evaluation also measured the widths of the welds.

As shown in figure 4, the EB weld had a uniform width of about 2 mm, and the

laser weld varied from 2 to 3 mm in width.

The solidification structures of the two welds were examined at higher

magnification. Figure 5 illustrates the fine columnar nature of the EB

weld. The columnar grain of the laser weld in figure 6 is similar in size,

but the higher travel speed resulted in less orientation in the direction of

welding and so a different appearance.

A previous study indicated that inclusion content reduces toughness. 3 Unetched micrographs of the welds (figures 7 and 8) show a higher inclusion

content in the laser weld, another reason for its lower toughness. The

laser weld had 1.7 x 10 4 inclusions per mm2

; the EB weld had 1.2 x 10 4 inclusions per mm2

.

The CT fracture surfaces (figures 9 and 10) reveal a fully ductile

fracture mode, which is characteristic of austenitic stainless steel at

4 K. The ductile dimples initiate at inclusions, so the dimple density is

another measure of the inclusion density. Fracture - surface dimple densities

for these welds and more conventional welds are given in table 8. The lower

dimple densities on the EB and laser welds correlate with their higher

toughness

.

Ferrite

Magnetic measurements were performed on the base metal and both welds.

All measurements on base metal, whether on the surface or cross sections,

indicated a ferrite number (FN) of zero (fully austenitic). The laser weld

was measured at 2 FN on both unmachined surfaces and on a carefully polished

cross section. The EB weld was measured at 2 to 3 FN on both unmachined

surfaces and on a carefully polished cross section. The presence of the

magnetic delta- ferrite phase in the welds is attributed to the reduction

in austenitizing potential through nitrogen loss during welding and the

increased solidification rate.

CONCLUSIONS

1. The electron-beam and laser welding processes produced weld strengths

(near 950 MPa) that exceeded the strength of type 316LN stainless steel

at 4 K.

2. The welds exhibited toughness values that averaged 30 percent larger

than those of more conventional welds of comparable strength.

3. The good toughness was attributed to the lower inclusion content in

these high-energy-beam welds.

304
ACKNOWLEDGMENTS

This work was partially supported by the Office of Fusion Energy, U.S.

Department of Energy. The authors thank C. N. McCowan for assistance in the

4-K tensile testing and P. T. Purtscher for assistance in the 4-K fracture

testing. The EB weld was produced by Sciaky and the laser weld was produced

by Combustion Engineering.

REFERENCES

1. H. I. McHenry, 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 (1985), pp . 121—122.

2. R. L. Tobler, T. A. Siewert, and H. I. McHenry, Strength—toughness

relationship for austenitic stainless steel welds at 4 K, Cryogenics

26, July (1986), pp. 392-295.

3. T. A. Siewert and C. N. McCowan, The role of inclusions in the fracture

of austenitic stainless steel welds at 4 K, in "Proceedings of International Symposium on Welding Metallurgy of Structural Steels," The

Metallurgical Society, Warrendale, Pennsylvania; in press.

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

nitrogen in stainless steel SMA welds for cryogenic service, Welding

Journal 66, March (1987), pp. 84-s—92-s.

5. N. J. Simon and R. P. Reed, Strength and toughness of A1SI 304 and 316

at 4 K, in "Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures—IX," NBSIR 86-3050, National Bureau of Standards,

Boulder, Colorado (1986), pp. 27—42.

6. J. A. Brooks, Weldability of high N, high Mn austenitic stainless

steel, Welding Journal 54, June (1975), pp. 189-s—195-s.

7. C. M. Weber, New developments in electron beam welding, in "Proceedings, Trends in Welding Research in the United States, S. A. David,

ed. , American Society for Metals, Metals Park, Ohio (1982), pp

647-652.

8. "Technical Data of Electron Beam Welding for Austenitic Stainless

Steel—Sciaky Machine EL 10," Mitsubishi Heavy Industries, Kobe

Shipyard and Engine Works, Kobe, Japan (1985).

9. E. A. Metzbower, D. W. Moon, and F. W. Fraser, Mechanical behavior of

laser beam weldments, in "Proceedings, Trends in Welding Research In

the United States," S. A. David, ed. , American Society for Metals,

Metals Park, Ohio (1982), pp . 563-580.

305
10. A. Nyilas, Laser beam welding of advanced superconducting cable

conduits, in "Proceedings of 14th SOFT," Avignon (1986); in press.

11. R. P. Reed, N. J. Simon, P. T. Purtscher, 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," NBSIR 86-3050, National Bureau of Standards,

Boulder, Colorado (1986), pp . 15—26.

12. R. L. Tobler, D. T. Read, and R. P. Reed, Strength/toughness relationship for interstitially strengthened AISI 304 stainless steel at 4 K

temperature, in "Fracture Mechanics: Thirteenth Conference," ASTM STP

743, R. Roberts, Ed., American Society for Testing and Materials,

Philadelphia (1981), pp . 250—268.

13. 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.

14. T. A. Siewert, Predicting the toughness of SMA austenitic stainless

steel welds at 77 K, Welding Journal 65., March (1966), pp . 23—28.

306
Table 1. Strength, Toughness, and Quality-Index Data Developed

on This and Other Heats of 316LN Base Metal at 4 K

11

Yield Strength,

ay

(MPa)

Fracture Toughness,

K

Ic

(J

,

) (MPa»m^)

QI,

K

Ic

(J) 7y

(MPa 2 • m’2

)

855-950a 185 1.58-1.76 x 10 5 935 140 1.31 x 10 5 1080 220 2.38 x 10 s 830 360 2.99 x 10 5 1060 150 1.59 x 10 5 1010 210 2.12 x 10 5 710 250 1.78 x 10 5 550 400 2.20 x 10 5

3

Heat used in this study. The range of yield strength was attributed to nitrogen segregation.

Table 2. Composition (in wt.%) of Type 316LN

Stainless Steel Plate Used in This Study

Element Heat Analysis NBS Analysis

C 0.016 —

Mn 1.10 —

Si 0.54 —

P 0.020 —

S 0.015 —

Cr 17.97 —

Ni 10.51 —

Mo 2.12 —

N 0.14 0.15

0 — 0.009

Cu 0.17 —

Co 0.25 —

307
Table 3. Welding Parameters

Electron Beam

Accelerating voltage 50 kV

Beam current 175 mA

Beam oscillation frequency 60 Hz

Beam oscillation width 2 mm Welding speed 200 mm min

Focal plane 38 mm in front of part

Gun-to-work distance 152 mm Weld direction vertical up (horizontal gun)

Laser

Beam power 8 kW (at workpiece)

Gas shield 450 1/hr He^ (3-flow nozzle)

Welding speed 500 mm min

Penetration 14 mm/ side in 2-pass /2-side butt weld

Focal length 375 mm using F5 parabolic mirror

Weld position flat

Table 4. Transverse Tensile Data at 4 K

Specimen EB Specimen Laser

EB-1 EB-2 Average Laser-1 Laser-2 Average

Yield strength (MPa) 981 943 962 940 940 940

Tensile strength (MPa) 1220 1190 1205 1130 1080 1105

Elongation (%)

a 9.4 8.7 9 b b b

Reduction in area (%) 14.6 16.3 15.4 b b b

Fracture location c c — d d —

a. Lower values than typical for base metals because the specimens were oriented in the short transverse direction.

b. Not available.

c. 3 mm from fusion line near tensile specimen shoulder.

d. 12 mm from fusion line near tensile specimen shoulder.

308
Table 5. Charpy V-notch Data at 76 K

Weld Electron Beam Laser

Notch direction

(with reference to

welding direction)

With Against With Against

Absorbed energy (J) 81 81 63 68

78 84 71 63

Average 80 82 67 66

Average by process 81 66

Table 6. Charpy V-notch Data Developed Previously

for Conventional Welding Processes at 77 K

Weld

Process

Alloy

Type

CVN

Absorbed Energy

(J)

SMA 316 24

SMA 308 16

SMA 308L 29

SMA 308L 22

SMA 308L 26

SMA 308L 27

SMA 308L 16

SMA 308L 34

SMA 308 22

SMA 316 42

SMA 316L 45

SMA 308L 24

SMA 308L 34

309
Table 7. Fracture-Toughness and Quality-Index Data for EB and Laser Welds

Specimen Notch

Direction 3 Jlc

- 2 (kJ.m )

K

Ic

(J > (MPa.m 1

)

QI,

Kt

( J )

• a Ic , y

(MPa*m^)

EB , CT-1 against 118 157 1.51 x 10 5 EB , CT-2 with 117 156 1.50 x 10 5 Laser , CT-1 against 104 147 1.38 x 10 s Laser , CT-2 with 101 145 1.36 x 10 5

With or against the direction of welding

Table 8. Fracture-Toughness and Quality-Index Data for Some Previously

Tested Welds with Yield Strengths above 850 MPa 2

Weld

Process

Alloy

Type (mp!)

K

Ic

(J

l

(MPa-m"*)

Q1 ^ (MPa. m

2 ) SMAW 316L 879 95 8.35 x 10

4 SMAW 316L 886 108 9.57 x 10*

SMAW 316L 881 98 8.63 x 10

4 SMAW ' 316 924 154 1.42 x 10

s FCAW 308L 1117 66 7.37 x 10

4 GMAW 308L 918 132 1.21 x 10

5 SAW 316L 860 99 8.51 x 10

4 SMAW 316L 882 171 1.51 x 10

s

Table 9. Fracture-Surface Dimple Densities for Various Welds

Weld

Process

Alloy

Type

Dimple Density

(mm

2

)

EB

3

316LN 1.7 x 10|

Laser 316LN 2.15 x 10

5 SMAW 308LN 3.2 x 10

5 SMAW 308LN 3.5 x 10 -

This study

310
If

YIELD STRENGTH, ksl

0 50 100 150 200 250

0

YIELD STRENGTH, MPa

Figure 1. Strength—toughness relationships for AISI type 316LN austenitic

stainless steel and matching composition welds.

Figure 2. Butt-joint weld configuration.

311

FRACTURE

TOUGHNESS,

ksiVF
weld centerline

C-3

Figure 3. Orientation of mechanical test specimens removed from weld.

Figure 4. Polished cross sections of laser (left) and EB (right) welds.

312
Figure 5

Figure 6.

Micrograph of the EB fusion line showing the fine columnar

dendrites in the weld (on the right). Orientation: transverse

to the welding direction.

Micrograph of the laser fusion line showing the coarse columnar

dendrites in the weld (on the right). Orientation: transverse

to the welding direction.

313
Figure 7. Higher magnification micrograph of unetched EB weld showing very

few pores or inclusions.

Figure 8

.

Higher magnification micrograph of unetched laser weld showing a large number of pores and inclusions.

314
Figure 9. Micrograph of the CT fracture surface of the EB weld.

Figure 10. Micrograph of the CT fracture surface of the laser weld.

315

=l#^=S =Jt

M mj- ¥MHFt r Mt-=^==i A

ff %r ft #» «*»#'#

# Jt #» Mi %

r >

i

1

B

I

317

HORMETALUCS PROGRAM

CONTRIBUTORS: R. D. Kriz and M. B. Kasen

OBJECTIVES

• Development of improved methods for screening the influence of component

parameters on composite material performance

• Application of these methods to select optimal components for radiationresistant, organic-matrix insulators to be used in superconducting magnets

of magnetic fusion energy systems

• Development of analytical methods to predict the influence of component

properties and internal stress state on damage development in composite

materials and to predict the influence of such damage on the mechanical

performance of composite materials

• Application of these techniques and methods in the development of improved

electrical insulating materials for superconducting magnets subjected to

neutron irradiation

RESEARCH HIGHLIGHTS

• An integrated system was developed for efficient, low-cost specimen

composite production and mechanical performance screening. The 3.2-mm

diameter rod specimens are produced in-house under excellent control.

The rods need only be cut to length before testing. Procedures and

facilities have been developed for testing the flexural, short-beam

shear, and torsional moduli and the fracture strength, G^ , of the rods

at temperatures from 295 K to 4 K.

• Ninety-eight unreinforced resin specimens, an equal number of specimens

reinforced with 48-volume-percent type-E glass fabricated from three epoxy

systems, and one bismaleimide system have been submitted to Oak Ridge

National Laboratory for irradiation at 4.2 K in the National Low Temperature Neutron Irradiation Facility. The specimens will be irradiated with

three increasing doses, warmed to room temperature, and returned to NBS

for testing.

• Laminated plate analysis was used to demonstrate how optimal component/;

could be fabricated from knowledge of component properties and other

laminate design variables. A micromechanical model of a graphite-epox v

fiber-reinforced laminate was used to demonstrate the relationship between

thermal contraction and local residual stress due to cooling to ervogon:

temperatures. Analysis also showed that replacement of the Inconel memr^r

of a copper-inconel composite with beryllium copper would signif icantlv

lower the thermally induced residual stresses.

319



LAMINATED PLATE ANALYSIS FOR LOW TEMPERATURE APPLICATIONS:

METALLIC AND NONMETALLIC LAMINATES

R. D. Kriz

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

The theory for laminated plate analysis is outlined. Simple

examples illustrate mechanical load and thermal stress analysis

in the temperature range 76 to 300 K for laminates reinforced with

metallic and nonmetallic materials. Laminates studied were copper

reinforced with Inconel, copper reinforced with a Cu—Be alloy,

and a fiber-reinforced graphite—epoxy crossply laminate, [0/90] s

. Although layers of Inconel provided good mechanical reinforcement

for layers of copper, the Cu—Be alloy provided the same reinforcement with lower residual thermal stresses in the copper. For

designing laminates with a specific coefficient of thermal ex- pansion, nonmetallic fiber-reinforced laminates are preferred.

At 76 K, moisture in polymeric composite laminates eliminates

thermally induced stresses.

INTRODUCTION

The design of fusion energy structures includes laminated material

systems that can either be metallic or nonmetallic and isotropic or anisotropic. Nonmetallic anisotropic laminated materials are usually selected

for structural applications because the individual layers of fiber

reinforced materials can be chosen to obtain specific properties for the

laminate. 1 Although less common, isotropic metals can also be used to

reinforce other metals in a laminate configuration. Components of metallic

laminates synergistically improve performance; for example, the Japanese

samurai swords derived their strength from their metallic laminate

composition. 2

The toroidal field (TF) coils recently designed for the compact

ignition tokamak (CIT) include copper conductor reinforced with Inconel.

In this application, the load is applied normal to the laminate midp lane.

Studies by Becker et al. 3 demonstrated that reinforcing copper with Inconel

significantly increases its compressive yield strength.

321
This improved compressive property was measured at the National Bureau

of Standards on full-scale laminates that were created by explosively bonding copper and Inconel plates. 4 Measurements focused on the copper—Inconel

interface bond. Without a good interfacial bond between layers, reinforcement is not possible and many structural advantages are sacrificed. The

interfacial bond strength is especially critical in laminates constructed of

highly anisotropic layers because differences in Poisson's contractions can

cause large strain gradients across the interface. For metallic laminates,

variations in Poisson's ratio among the layers are not large, but differences in thermal contraction (over the temperature range 4 to 300 K) can cause large stress gradients near the interfaces. When thermal strains are superposed with strains induced by a mechanical load, these additional

interfacial thermal strains can be a significant factor in the design of

metal structures for low temperature applications.

This study was undertaken to provide designers of fusion energy structures with an outline to design laminates constructed with metallic and

nonmetallic materials that may be suitable in the temperature range 4 to

300 K. In particular, we give examples of nonmetallic graphite—epoxy laminates and the metallic laminates similar to materials used in the TF coils

of the CIT. We also illustrate how the coefficients of thermal expansion

can be predicted from laminated plate theory.* Hence, structural components

can be designed to operate at low temperatures with minimal problems in

thermal mismatch with other structural members.

LAMINATED PLATE THEORY

Comprehensive development of laminated plate theory is found in references 1 and 5. Using the same definitions and notation of these references,

this report presents an overview of the theory that is relevant to design

parameters

.

Laminated plate theory is based on the assumption that adjacent layers

are perfectly bonded and in a state of plane stress (ct z = 0) normal to the

thin x—y plane of the laminate. The theory also adopts the Kirchhof f—Love

assumptions for bending of thin plates:

1 . Normals to the laminate midplane remain straight and normal to

the deformed midplane. Hence, out-of -plane shear deformations

(r xz - r

y

Z - 0) are assumed to be zero.

*The computer program used to predict the properties for these examples is

not described fully in this report. This NBS computer program is public

domain and can be obtained by sending a formatted, blank, 360-kilobyte,

5-1/4-in floppy disk to R. D. Kriz, MC 430, National Bureau of Standards,

Boulder, CO 80303. The program is designed to operate on any personal

computer that uses DOS 3.0.

322
2. The length of midplane normals does not change. Hence, outof-plane deflections, w, are associated only with in-plane

stretching: w = w(x,y).

We also assume small deflections for thin, elastic plates.

Mechanical Loads

With these assumptions, stresses ( crx , cry, r xy) at a point can be

calculated for each layer (k) from the plate curvatures («x , Ky, «Xy) anc* in-plane strains (e£, e y> 7Xy>) at the midplane.

CT X Qii Q

2 Q i

r r 0 > e x

r > *x

* ay

=

Ql 2 Q22 Q26 < * 0

e y

* -zk- «y (1)

r xy

w j k

. Q 1 6 Q26 Qe6

k

0

w 7xy . - *xy

J

where are the elastic properties defined in reference 5, and Zk is

defined m Fig. 1. The force-moment resultants are defined in Fig. 2.

rt/2 rt/2 pt/2

Nx " ax

dz

N

y " CTydZ , NXy

= T XydZ

• -t/2 -t/2 • -t/2

r t/2 pt/2

r t/2 = ^^zdz j %

- yZ dz

MXy

= T ^yZciz

« -t/2 * -t/2 * -t/2

For discrete layers, the integration in Eq. (2) is summed over N layers:

N 1

: iN x N N

c r 'v ffx Ny 2 * Oy

NXy

k = 1 ^k -

r xy

Mn 1 x n

* zk ^x

My 2 Oy

MXy

J k = 1 Zv. 1

. r xy «

By substituting Eq. (1) into Eq. (3) and assuming that each layer is

homogeneous , we obtain

323
f N 1 N

' Qn Ql 2 Ql 6

r e° 1 x

r zk

N

y

2 Ql 2 Q22 Q2 6

* 0

e y

dz +

> Nxy - k = 1

. Ql 6 Q26 Q6 6 . k

0

. "Yxy - • Zk -

1

' Ql! Qi 2 Ql 6

f >

/c x

pZk

+ Ql 2 Q22 Q2 6

•< Ky zdz

. Ql 6 Q2 6 Q6 6 . k

< *xy - • Zk -

1

f nM 1 x N

' Qi 1 Qi 2 Ql 6

f e°e 1 x My 2 Qi 2 Q22 Q2 6

4 0

€ y k - 1 — — — 0 MXy «. j . Ql 6 Q2 6 Qe e . k

k T^xy j •

Zk

Zkzdz

1

+

+

_ r *\ pZk

> Qn Qi 2 Ql 6 «x

Qi 2 Q2 2 Q26 4 Ky z

2 dz *

. Ql 6 Q26 Q6 6 . k

- Kxy - • Zk1

j

( 4 )

(5)

Combining like terms in Eqs. (4) and (5), integrating, and rearranging the

variables yields one relationship with partitioned submatrices:

Nx

’ A 1

1

^12 ^16 Bn Bi 2 ®1 6

f e° 1 x N

y

A 1 2 A2 2 A26 Bj 2 B 2 2 &2 6

c 0y

Nxy A 1 6 A2 6

A. 66 6 B 2 6 Be 6

f °

e xy

4 < Mx

B 1 1 Bi 2 ® 1

6

Du Di 2 Die AC X

M

y

B 1 2 B 2 2 ®26 D x 2 D 2 2 D 2 6 /Cy

MXy

B 1 6 ®2 6 B 66 Dl 6 D 2 6 Dee KXy

N

where Aij 2 (Qij (z^ — zk-l)

k 1

N — 2 2 B ij “ 2 (Qij ) ^

( zk

— Zk -

1

)/2,

k - 1

( 6 )

324
and 'ij

N

2

k - 1 (Qij h

(Zfc ” 4-i)/3

Equation (6) is written in the form where the midplane strains and

curvatures are prescribed and the midplane stress and moment resultants

are calculated for a particular laminate. Other forms of Eq. (6) can be

derived; 1 there are four possible variations. For example, in one variation, the moment resultants (Mx , My, Mxy) and stress resultants (Nx , Ny,

NXy) are prescribed and the corresponding strains (e£, A£y) and curvatures (/cx , Ky, Kyy) at the midplane are calculated. Of the four possible

combinations, three are expressed in abbreviated notation in Eqs . (7)

through (9)

N

M

(7)

e o as A" 1: —a' 1 b N

M _ BA-1 :d - ba_1 b K

^ ° \ J

( 8 )

4 £ o =» A* - B* D*

K _D*-i H

*

L j

where A = A' B* - -A’ 1 B, H*

H B D

n* :

i

N < M

^ j

(9)

BA-1

, and D = D — BA _1 B. Stresses in each

layer can now be calculated from the following relationship.

{o } k = [Q]k (U 0

) - Zk {/c}} ( 10 )

Thermal Loads

At low temperatures, additional strains are the result of thermal

expansion and contraction, but free thermal strains alone do not cause

stress in homogeneous materials. When laminates are cooled to low temperatures, the only source of stress is the thermal mismatch of different

layers, which causes constraint of adjacent layers. This mismatch can be

modeled by laminated plate analysis. In each layer (k) , the total strain

({e} k) is the sum of the mechanical strains ((e CT

) k) and the thermal strains

(UT

>k>-

(Ok - U

CT >k + (* T

)k

In a comprehensive analysis, strains caused by other environmental effects

are included in Eq. (11). For example, absorbed moisture or irradiation can cause swelling of polymer-base laminates. However, in this outline we cot. sider only mechanical strains and thermal strains.

325
The strains in Eq. (1) are the strains that cause stresses ({e*7 }^).

Hence, to include thermal strains in the laminate analysis, we solve for the

stress - induced strains in Eq. (11) and substitute them into Eq. (2). In

schematic form, Eq. (2) is rewritten in terms of total strains ((e)) and

thermal strains ({e^}):

(N> =

r h/2

{u} k dz =

•'-h/2

(M) =

, h/2

(a) k zdz =

•'-h/2

rh/2

[Q] k “ fT )k dz

•'-h/2

rh/2

[Qlk ~ « T

)k zdz

—h/2

( 12 )

For each layer (k)

,

the thermal strains in Eq . (12) can be rewritten in

terms of a temperature change (AT) and the coefficients of thermal expansion ((a)):

U

T

} k = AT{a} k (13)

Substituting Eq . (13) into Eq . (12) and integrating yields

< N =» A B

M

k. j B D

where

(NT

) - AT

r V2

(Qlk l

• -h/2

and

(MT

) - AT

p h/2

[Q] k (

*1 -h/2

(14)

are the midplane thermal loads (force and moment resultants).

The laminate coefficients of thermal expansion are calculated from

Eq. (14) for the special case where the curvatures ({/c}) are zero, the

mechanical loads ({N} and (M)) are zero, and the laminate is symmetric

([B] = 0). Hence, Eq . (14) decouples into

[A] {

e

0

} =• {NT

} (15)

326
Solving Eq. (15) for midplane strain ( { e 0

) ) yields

fh/2 _ U°) = [A]’ 1 {NT

} = AT [ A ]

‘ 1 [Q] k {a} k

dz

(lg)

J-h/2

For the laminate, we can define this relationship:

{ eaT

) = AT{a) (17)

where {a} are the laminate coefficients of thermal expansion. When only

thermal strains are present, {e 0

} = {e 0 ^}, and by inspection of Eqs . (16)

and (17), we arrive at an expression for {a}:

(a) - [A}* 1 ph/2 _

[Q] k

dz

J-h/2

(18)

Now the stresses in each layer can be calculated for a variety of mechanical

and thermal load conditions:

{a} k = [Qlk l U°) + zk («) “ {<*}kAT > (19)

Stress Approximations for the Laminated Edge

In the previous plate analysis, plane stress is assumed through the

thickness of the thin plate; therefore, the interlaminar stresses (ct z , r xz , ryZ ) through the thickness are negligible. For interlaminar stresses in

thick plates or for edge stresses in thin plates, the plane stress approximation is no longer valid. For thick plates, interlaminar stresses can be

approximated by modifying thin laminate plate theory, 6 ’ 7 but a finite element model is a more accurate approach. 8 For thin laminated plates, these

stresses exist only near the plate boundaries (edges) , where stresses normal

to the edge surface are zero. Originally, solutions for these edge stresses

were obtained numerically; 9 more recently, complicated, exact solutions have

been obtained. 10 ’ 11

For simplicity, we use an approximation proposed by Pagano and Pipes. 12 The actual and approximate distributions for the interlaminar edge stresses

(<7 Z ) are shown in Fig. 3. A relationship for the amplitude (a ra ) is obtained

by summing the moments about the x-axis for the k layers shown in Fig. 1

»h oyt d£ (20)

zk

Unfortunately, similar approximations cannot be made for the shear stress'-;

(r xz . r yz)

Thus, for each layer, we approximate only the interlaminar

peeling stress (cr z ) k by summing stresses in the y direction (oy) for each

layer (k)

In the computer program, we integrate over four evenly spaced

points for each layer.

45

a

Z. " ani - 14h 2 k

327
RESULTS AND DISCUSSION

With this analysis included in a computer program, we studied three

simple laminate configurations: Inconel/Cu/Inconel ; Cu—Be/Cu/Cu—Be ; and

graphite—epoxy [0/90] s

. Computer program input data files with results

for each of these configurations are given in Appendices A, B, and C,

respectively

.

Incone 1/Cu/Inconel

Inconel/Cu/Inconel , an isotropic metallic laminate, was chosen as a simple example to illustrate the stresses caused by thermal mismatch among

the layers. Layer thicknesses are the same as those of laminates tested in

reference 3. The material properties and dimensions used by the computer

program are shown in Fig. A.l as a data file. In this data file,

o EL is Young's modulus in the textured (L) direction

(defined in Fig. 1);

o ET is Young's modulus transverse to the L-direction;

o GLT and NULT are the shear modulus and Poisson's ratio

in the L—T plane;

o TEXPL and TEXPT are the thermal coefficients of expansion

in the L- and T-directions

;

o MEXPL, MEXPT, and CV are moisture properties that apply

only to swelling of polymeric materials.

o NX is the laminate load applied in the x-direction

(NX and other loads are defined in Fig. 2); the angle

between the textured L-direction and the x-axis is

defined for each layer.

In this example, no texture exists; hence, EL - ET and TEXPL - TEXPT.

(However, metal textures, even small ones, can be included in the analysis.)

Also, NX — NY — NXY - MX — MY - MXY - 0 because no loads were applied in

the L—T plane of the laminate. When thermal loads are studied, we assume

a stress-strain-free state at room temperature (300 K) and prescribe an operating temperature of 76 K.

For this problem, only a portion of the results from the computer

program are shown in Fig. A. 2. Stresses are presented in a table and a graph; the graph represents only stress components only in the x, y, and z directions

.

Since there is no texture in this example, all shear stresses are zero,

and ct x (SIG X) equals Oy (SIG Y) . For operating temperatures below room

temperature, the opposite is true. Most significant is the large tension

stress of 117 MPa (17 ksi) in the copper at 76 K, which is 42% of the yield

strength of copper. Thus, for this laminate, yielding occurs at much lower

mechanical loads at 76 K than at room temperature.

328
Cu—Be /Cu/Cu—Be

For comparison, we consider a laminate in which a Cu—Be alloy replaces

the Inconel reinforcement in the first example (see Fig. 5.4.2-El of reference 13). The input data file is shown in Fig. B.l, and the results are shown in Fig. B.2.

The most significant result is that the residual thermal stress in the

copper changes to a compressive stress of —96.5 MPa (—14.8 ksi) ; hence, much

larger mechanical loads can be experienced before yielding occurs in the

copper layer. Unfortunately, the interlaminar stress (cr z ) at the laminate

edge changes from compressive to tensile (36 MPa; 5.3 ksi). If poor bonding exists between the Cu and Cu-Be layers, delaminations can occur if tensile loads are applied in the x—y plane of the laminate. If compressive

loads are applied in the z-direction, these interlaminar stresses become

insignificant

From these results, we conclude that a close match of the coefficient

of thermal expansion among layers results in low thermal stresses. Hence,

a good reinforcement does not always result in a poor thermal mismatch and

high residual thermal stresses.

The coefficient of thermal expansion for the Cu—Be/Cu/Cu—Be laminate

was 16.5 10" 6 /K and for the Inconel/Cu/Inconel laminate, 11.4 10" 6 /K.

Therefore, the designer must allow for significant residual thermal stresses

when a particular coefficient of expansion is required by the designer. For

metals, thermal mismatch is mainly due to differences in thermal expansion

among the layers, seldom to differences in their Poisson's ratios. For

laminates with fiber-reinforced composite layers, differences in Poisson's

ratios are often significant.

These two examples demonstrate that metal laminates offer less control

of design variables than composite-fiber-reinforced laminates, which have a much wider range of design variables (anisotropic material properties) that

can be controlled by laminate theory.

Graphite—Epoxv f 0/901 ,-.

This graphite—epoxy laminate, with a [0/90] s stacking sequence, was subjected to three different loads: (1) mechanical load only, (2) thermal

load only, and (3) moisture load only. The data file is given in Fig. C.l

and results are shown in Figs. C.2 through C.4. For ease of comparison, all

stresses are plotted on the same scale.

For this laminate with a moderate load of 4.2 kN/cm (2900 Ibf/in)

applied in the x-direction, the stresses that are calculated are shown in

Fig. C.2. Because of elastic anisotropy, the stresses in the x- and ydirections are significantly different. The interlaminar edge stresses are very low. Shear stresses are, of course, negligible for laminates with 0

and 90° fiber orientations, but laminates with ±45° layers have much larger

shear stresses.

329
REFERENCES

1. R. M. Jones, Mechanics of Composite Materials . McGraw-Hill, New York,

1975; pp. 147-237.

2. B. Dodd, "The Making of Old Japanese Swords," Journal of Mechanical

Working Technology vol. 2, 1978; pp . 75—84.

3. H. Becker, M, Besen, E. Bobrov, J. Chen, and D. B. Montgomery,

"Structural Properties of Reinforced Copper," Internal report PFC/RR86-12, Plasma Fusion Center, Massachusetts Institute of Technology,

Cambridge, Massachusetts, 1986.

4. Graham Brown, Plasma Physics Laboratory, Princeton University,

Princeton, New Jersey, private communication, 1987.

5. J. E. Ashton and J. M. Whitney, Theory of Laminated Plates . Technomic

Publishing Co., Stanford, Connecticut, 1970.

6. J. M. Whitney, The effect of transverse shear deformation on the

bonding of laminated plates, Journal of Composite Materials , vol. 3,

1969; p. 534.

7. J. M. Whitney and N. J. Pagano , Shear deformation in heterogeneous

anisotropic plates, Journal of Applied Mechanics , vol. 37, 1970; pp

.

1031-1036.

8. R. D. Kriz, Influence of ply cracks on fracture strength of graphite/

epoxy laminates at 76 K, Effects of Defects in Composite Materials . ASTM STP 836, American Society for Testing and Materials, Philadelphia,

1984; pp. 250-265.

9. R. B. Pipes and N. J. Pagano, Interlaminar stresses in composite

laminates under uniform axial extension, Journal of Composite

Materials, vol. 4, 1970; pp . 538-548.

10. S. S. Wang and I. Choi, Boundary layer thermal stresses in angle-ply

composite laminates, Modern Developments in Composite Materials and

Structures . J. L. Vinson, ed. , ASME, New York, 1979; pp . 315-341.

11. R. I. Zwiers, T. C. T. Ting, and R. L.Spilker, On the logarithmic

singularity of free-edge stress in laminated composites under uniform

extension, Journal of Applied Mechanics , vol. 49, 1982; pp . 561-569.

12. N. J. Pagano and B. P. Pipes, Some observations on the interlaminar

strength of composite laminates, International Journal of Mechanical

Science . vol. 15, 1973; pp. 679-688.

13. Handbook for Superconducting Machinery . Metals and Ceramics Information

Center, Battelle Columbus Laboratories, Columbus, Ohio, 1977.

330
The results for thermal loads only are shown in Fig. C.3. For this

laminate, stress-free states exist above room temperature (300 K, 81°F).

For comparison, the same operating temperature (76 K, —323°F) was used for

all three laminates. Under thermal load, the interlaminar edge stresses

were three times larger than under mechanical load. With a combination of

thermal and mechanical loads, delaminations are most likely to occur between

90° layers; the thermal and mechanical stresses are nearly identical. Compared with that of the metallic laminate in the second example, the thermal

expansion is quite low. With a parametric study, it is possible to obtain a stacking sequence that results in zero thermal expansion, but large residual

thermal stresses are unavoidable.

Residual stress states caused by absorbed moisture are shown in

Fig. C.4. When moisture is absorbed by the epoxy matrix, each layer swells

only in the T-direction. This swelling produces an effect which is opposite that of the thermal load; hence, all residual stresses are opposite

in sign. At 1.0% moisture weight gain in the laminate, these moistureinduced stresses can cancel the thermal stresses, leaving this laminate in

a stress-free state at 76 K.

CONCLUSIONS

Design methods deriving from laminate plate theory can be applied

advantageously to both metallic and nonmetallic laminates. For nonmetallic

laminates there are a wide variety of material properties that can be used

in a laminate design. Although there is less choice among the properties

of metals, the properties of a specific layer can be improved by adding

reinforcement layers.

To minimize residual thermal stresses in low temperature applications,

all layers within the laminate must have closely matching coefficients of

thermal expansion. Metallic laminates are best suited for this purpose.

In this study, both Cu—Be and Inconel provided good reinforcement, and the

Cu—Be reinforcement resulted in low thermal stresses.

When a laminate is designed for a specific coefficient of thermal

expansion, a large difference in the coefficients of thermal expansion

among the layers is required. Nonmetallic materials are best suited for

achieving a specific coefficient of thermal expansion for the laminate but

large, thermally induced residual stresses are created. At 76 K, moisturein polymeric composite laminates can eliminate these thermally induced

residual stresses.

Although these concepts are easily understood, the laminate plate

theory outlined in this report is a tool for quantitative laminate design

when a specific material response is required in a low temperature

application.

331
Fig. 1 Laminate configuration defined.

Fig. 2 Midplane loads defined.

332
Fig. 3

(+)

<*z.

(-)

Actual and approximate interlaminar stress

distributions defined for the o z stress.

333
Appendix A. Computer Data File and Results for a Copper Laminate

Reinforced with Inconel

*** COPPER /REINFORCED-INCONEL LAMINATE THERMAL LOADS, R. D. KRIZ 12/3/86 ***

NO. OF LAYERS- 3 NUMBER OF MATERIAL TYPES- 2

LAYER NO. 1 ANGLE-+00.0 THICKNESS—1-0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -INCON

LAYER NO. 2 ANGLE—i-OO.O THICKNESS—1-0 . 0100 MATERIAL TYPE- 2 PLOT TYPE - CU

LAYER NO. 3 ANGLE—+00 . 0 THICKNESS—+0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -INCON

MTL. TYPE 1 EL—1-0 . 300E+08 ET-+0 . 300E+08 GLT—1-0 . 155E+08NULT—+0 . 300E+00

TEXPL—(-1 . 050E-05 TEXPT—fl . 050E- 05 MEXPL—+0 . 000E+00 MEXPT-+0 . 000E+00

MOISTURE THRESHOLD VOLUME, CV- 0.0

MTL. TYPE 2 EL-+1 . 910E+07 ET-+1 . 910E+07 GLT-+7 . 350E+06NULT-+0 . 300E+00

TEXPL-+1 .412E-05 TEXPT-+1 . 412E-05 MEXPL—+0 . 000E+00 MEXPT-+0 . 000E+00

MOISTURE THRESHOLD VOLUME, CV- 0.0

NUMBER OF LOAD CASES- 1

LOAD CASE NUMBER- 1

+1 PLUS ONE INDICATES GIVEN LOADS AND MOMENTS CALCULATE STRAINS AND CURVATURE

NX-+0 . 000E+00NY—1-0 OOOE+OONXY—t-O . OOOE+OOMX-+0 . 000E+00MY—1-0 . 000E+00MXY-+0 . 000E+00

STRESS FREE TEMP—1-300.0 OPERATING TEMP—1-76.00 PERCENT MOISTURE ABSORBED- 0.000

-1 INDICATES THAT MAX MIN LIMITS ARE CHOSEN FOR PLOTTING STRESSES

Fig. A.l Input data file for laminate computer program for a copper layer

laminate reinforced with Inconel. Materials properties were taken

from reference 13.

334
••* COPPER/REINFORCED-INCONEL LAMINATE THERMAL LOADS, R.D. KRIZ 12/3/86 •••

LOAD CASE NO.= 1

STRESS-FREE TEMP.= 3.00E+02 OPERATING TEMP.= 7.60E+01 MOIST. CONT.= 0.00E-01

MFCHANICALIMID PLANE LAMINATE LOADS

LOADING I NX= 0.00E-01 NY= 0.00E-01 NXY= 0.00E-01

MSW= 1 I MX= 0.00E-01 MY= 0.00E-01 MXY= 0.00E-01

IMIDPLANE LAMINATE STRAINS AND CURVATURES

I EXO= 0.00E-01 EYO= 0.00E-01 EXYO= 0.00E-01

PLY I KX= 0.00E-01 KY= 0.00E-01 KXY= 0.00E-01

ANGLE

PLY SIG L SIG T TAU LT SIG X SIG Y TAU XY

10 0

1 -8 . 39E+03 —8 . 39E+03 0. 00E-01 -8.39E+03 -8.39E+03 0.00E-01

20.0

2

1 . 68E+04 1 . 68E+04 0.00E-01 1 . 68E+04 1 68E+04 0.00E-01

30.0

3

-8 . 39E+03 —8 . 39E+03 0.00E-01 —8 . 39E+03 -8.39E+03 0.00E-01

SIGZ

0.00E-01

—3 . 75E+02

-1 . 50E+03

-3.37E+03

—5 . 99E+03

-8.24E+03

-8.99E+03

-8.24E+03

—5 . 99E+03

-3.37E+03

-1 . 50E+03

-3.75E+02

SYMBOL SYMBOL CORRESPONDING

PRIORITY PLOTTED STRESS

LOW 5 X —— STRESS IN X DIRECTION

4 — Y STRESS IN Y DIRECTION

3 S SHEAR STRESS IN X-Y PLANE

2 Z STRESS IN Z DIRECTION

HIGH 1 0 ZERO LINE

STRESS PLOT: LOAD CASE NO.= 1 MINIMUM STRESS= -0.89908E+04

MAXIMUM STRESS= 0.16783E+05

STRESS INCREMENT= 0.47000E+03

THICKNESS

INCREMENT

0 . 25000E-02

I PLY

TYPE

INCON

CU

INCON

MIN= -0 . 92000E+04

V

-STRESS- 0 TTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTT

H—Y 1 0 1 1 1

THICKNESS T ( I

LISTED DOWN

RIGHT

COLUMN

I

I

I ! 0 . 00E-01

1 9000E+05=MAX

V

! Y j 0 j J j ! 2 . 50E-03

! Y j Z 0 j j j ! 5 . 00E-03

! Y ! Z 0 { j j ! 7 . 50E-03

H—Y-

! Y ! Z 0 ! j ! ! 2 . 25E-02

! Y j Z 0 1 j

i 1 2 . 50E-02

! Y j 0 j j j ! 2 . 75E-02

-Y— ! 1

! 1

! 1

! 1

• ! 2

00E-02

25E-02

50E-02

75E-02

00E-02

-t—Y- - 0-

! 3 00E-02

COEFFICIENTS OF EXPANSION FOR LAMINATE X-Y COORD

X-DIRECT ION Y-DIRECTION

SHEAR XY-PLANE

THERMAL COEFF.

1 . 137E-05

1 . 137E-05

0.000E-01

MOISTURE COEFF

0 . 000E-01

0 . 000E-01

0 . 000E-01

Fig. A. 2 Output data file of results calculated by the laminate compnr.o r program

for data given in Fig. A.l. Stresses are given in tables and graphs.

335
Appendix B. Computer Data File and Results for a Copper Laminate

Reinforced with a Copper—Beryllium Alloy

*** COPPER/REINFORCED -COPPER: BERYLLIUM, THERMAL LOADS, R.D. KRIZ 12/3/86 ***

NO. OF LAYERS- 3 NUMBER OF MATERIAL TYPES- 2

LAYER NO 1 ANGLE—+00 . 0 THICKNESS—+0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -CU:BE

LAYER NO. 2 ANGLE—+00 . 0 THICKNESS—+0 .0100 MATERIAL TYPE- 2 PLOT TYPE - CU

LAYER NO. 3 ANGLE-+00.0 THICKNESS—+0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -CU:BE

MTL. TYPE 1 EL—+0 . 185E+08 ET-+0 . 185E+08 GLT-+0 . 712E+07NULT-+0 . 300E+00

TEXPL-+0 . 178E-04 TEXPT—+0 . 178E-04 MEXPL—+0 . 000E+00 MEXPT—+0 . 000E+00

MOISTURE THRESHOLD VOLUME, CV- 0.0

MTL. TYPE 2 EL-+1 . 910E+07 ET-+1 . 910E+07 GLT-+7 . 350E+06NULT-+0 . 300E+00

TEXPL-+1 . 412E-05 TEXPT—+1 .412E-05 MEXPL—+0 . 000E+00 MEXPT—+0 . 000E+00

MOISTURE THRESHOLD VOLUME, CV- 0.0

NUMBER OF LOAD CASES- 1

LOAD CASE NUMBER- 1

+1 PLUS ONE INDICATES GIVEN LOADS AND MOMENTS CALCULATE STRAINS AND CURVATURE

NX-+0 000E+00NY-+0 . 000E+00NXY-+0 . 000E+00MX-+0 . 000E+00MY-+0 . 000E+00MXY-+0 . 000E+00

STRESS FREE TEMP-+300.0 OPERATING TEMP-+76.00 PERCENT MOISTURE ABSORBED- 0.000

-1 INDICATES THAT MAX MIN LIMITS ARE CHOSEN FOR PLOTTING STRESSES

Fig. B.l Input data file for laminate computer program for a copper layer

laminate reinforced with a Cu—Be alloy. Material properties were

taken from reference 13.

336
• + + COPPER/RE I NFORCED-COPPER: BERYLLIUM. THERMAL LOADS. R.D. KRIZ 12/3/86

LOAD CASE NO . = 1

• • *

STRESS-FREE TEMP.= 3.00E+02 OPERATING TEMP.= 7.60E+01 MOIST. CONT.= 0.00E-01

MECHANICALIMID PLANE LAMINATE LOADS

LOADING I NX= 0.00E-01 NY= 0.00E-01 NXY= 0.00E-01

MSW= 1 I MX= 0.00E-01 MY= 0.00E-01 MXY= 0.00E-01

IMIDPLANE LAMINATE STRAINS AND CURVATURES

I EXO= 0.00E-01 EYO= 0.00E-01 EXYO= 0.00E-01

PLY I KX= 0.00E-01 KY= 0.00E-01 KXY= 0.00E-01

ANGLE

PLY SIG L SIG T TAU LT SIG X SIG Y TAU XY

10.0

1 7.42E+03 7.42E+03 0.00E-01 7.42E+03 7.42E+03 0.00E-01

20.0

2 -1 . 48E+04 -1 . 48E+04 0.00E-01 -1 .48E+04 -1 . 48E+04 0 , 00E-01

3

0.0

3

7 . 42E+03 7 . 42E+03 0.00E-01 7.42E+03 7.42E+03 0.00E-01

SIGZ

0.00E-01

3 31E+02

1 . 32E+03

2.98E+03

5.30E+03

7.28E+03

7.95E+03

7.28E+03

5.30E+03

2.98E+03

1 . 32E+03

3.31E+02

SYMBOL SYMBOL CORRESPONDING

PRIORITY PLOTTED STRESS

LOW 5 X —— STRESS IN X DIRECTION

4 Y STRESS IN Y DIRECTION

3 S SHEAR STRESS IN X-Y PLANE

2 Z STRESS IN Z DIRECTION

HIGH 1 0 ZERO LINE

STRESS PLOT: LOAD CASE NO.= 1 MINIMUM STRESS= -0. 14834E+05

MAXIMUM STRESS= 0.79470E+04

STRESS INCREMENT= 0.42000E+03

THICKNESS

INCREMENT

0 . 25000E-02

I PLY

TYPE

CU

MI N= -0. 1 7000E+05

V -STRESS- 0 . 82000 E+04=MAX

V

CU : BE

CU : BE

TTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTT

—I 1 1 1 1—0 1 Y0

THICKNESS T ( I

LISTED DOWN

RIGHT

COLUMN

I

I

I 00E-01

50E-03

00E-03

50E-03

00E-02

25E-02

50E-02

75E-02

00E-02

25E-02

50E-02

75E-02

00E-02

Y2

Y5

Y7 — ! 1 Z! 1

! 1 Z! 1 —Y2

Y2

Y2

Y2 —Y3

COEFFICIENTS OF EXPANSION FOR LAMINATE X-Y COORD

X—DIRECTION

Y-DIRECTION

SHEAR XY-PLANE

THERMAL COEFF.

1 . 655E-05

1 . 655E-05

0 . 000E-01

MOISTURE COEFF

0.000E-01

0 . 000E-01

0 000E-01

Fig. B .Output data file of results calculated by the laminate computer ;

for data given in Fig. A. 2. Stresses are given in tables and •. : ::

337
Appendix C. Computer Data File and Results for a Graphite/Epoxy

Fiber-Reinforced Composite Laminate with Two Layers

of 90° Material Reinforced by Two Layers of 0° material,

[0/90] s

*** GRAPHITE/EPOXY LAMINATE; LOADS AND STRESSES FOR (0/90)SYM. LAMINATE ***

NO. OF LAYERS- 4 NUMBER OF MATERIAL TYPES- 1

LAYER NO. 1 ANGLE—+00 . 0 THICKNESS—+0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -+00.0

LAYER NO. 2 ANGLE-+90.0 THICKNESS—+0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -+90.0

LAYER NO. 3 ANGLE-+90.0 THICKNESS-+0 .0100 MATERIAL TYPE- 1 PLOT TYPE -+90.0

LAYER NO. 4 ANGLE—+00 . 0 THICKNESS—+0 . 0100 MATERIAL TYPE- 1 PLOT TYPE -+00.0

MTL. TYPE 1 EL—+20 . 00E+06 ET-+2 . 100E+06 GLT-+0 . 850E+06NULT-+0 . 210E+00

TEXPL--0 . 230E-06 TEXPT—+14 . 90E-06 MEXPL—+0 . 000E+00 MEXPT—+5 . 560E-03

MOISTURE THRESHOLD VOLUME, CV- 0.0

NUMBER OF LOAD CASES- 3

LOAD CASE NUMBER- 1

+1 PLUS ONE INDICATES GIVEN LOADS AND MOMENTS CALCULATE STRAINS AND CURVATURE

NX-+2 . 400E+03NY—+0 . OOOE+OONXY-+O . 000E+00MX-+0 . 000E+00MY-+0 . 000E+00MXY-+0 . 000E+00

STRESS FREE TEMP-+000.0 OPERATING TEMP-+00.00 PERCENT MOISTURE ABSORBED- 0.000

+1 INDICATES THAT MAX MIN LIMITS ARE ASSIGNED BELOW FOR PLOTTING STRESSES MAX STRESS-+1 . 000E+04 MIN STRESS-- 1 . OOOE+04

LOAD CASE NUMBER- 2

+1 PLUS ONE INDICATES GIVEN LOADS AND MOMENTS CALCULATE STRAINS AND CURVATURE

NX-+0 . 000E+00NY—+0 . 000E+00NXY-+0 . 000E+00MX—+0 . 000E+00MY-+0 000E+00MXY-+0 . 000E+00

STRESS FREE TEMP-+100.0 OPERATING TEMP-- 323.0 PERCENT MOISTURE ABSORBED- 0.000

+1 INDICATES THAT MAX MIN LIMITS ARE ASSIGNED BELOW FOR PLOTTING STRESSES MAX STRESS—+1 . 000E+04 MIN STRESS-- 1 . 000E+04

LOAD CASE NUMBER- 3

+1 PLUS ONE INDICATES GIVEN LOADS AND MOMENTS CALCULATE STRAINS AND CURVATURE

NX-+0 . 000E+00NY—+0 . 000E+00NXY—+0 . 000E+00MX-+0 . 000E+00MY-+0 . 000E+00MXY-+0 . 000E+00

STRESS FREE TEMP-+000.0 OPERATING TEMP-+00 . 00 PERCENT MOISTURE ABSORBED-+1 . 000

+1 INDICATES THAT MAX MIN LIMITS ARE ASSIGNED BELOW FOR PLOTTING STRESSES MAX STRESS-+1 . OOOE+04 MIN STRESS-- 1 . 000E+04

Fig. C.l Output data file for laminate computer program for a symmetrical

crossply laminate with a stacking sequence of [0/90] . Material

properties were taken from reference 8.

338
Fig. C

•** GRAPHITE/EPOXY LAMINATE; LOADS AND STRESSES FOR (0/90)SYM. LAMINATE ***

LOAD CASE NO .= 1 STRESS-FREE TEMP.= 0.00E-01 OPERATING TEMP.= 0.00E-01 MOIST. CONT.= 0.00E-01

MECHANICAL I MID PLANE LAMINATE LOADS

LOAD I NG

MSW= 1

PLY

ANGLE

I NX= 2.40E+03 NY= 0.00E-01 NXY= 0.00E-01

I MX= 0.00E-01 MY= 0.00E-01 MXY= 0.00E-01

IMIDPLANE LAMINATE STRAINS AND CURVATURES

I EXO= 5.41 E—03 EYO=—2 . 1 6E—04 EXYC^-9 . 51 E-1

I KX=-3 . 09E—09 KY=-3 . 97E— 1 0 KXY= 2.30E-16

PLY SIG L SIG T TAU LT SIG X SIG Y TAU XY

10.0

1

1 . 09E+05 1 . 94E+03 -8.08E-05 1 . 09E+05 1 . 94E+03 —8 . 08E-05

2 90.0 - 2

-1 . 94E+03 1 . 13E+04 4 99E-04 1 . 13E+04 -1 . 94E+03 8 08E-05

3

90.0 - 3

-1 . 94E+03 1 . 13E+04 4.99E-04 1 . 13E+04 -1 . 94E+03 8.08E-05

40.0

4

1 . 09E+05 1 . 94E+03 -8 . 08E-05 1 . 09E+05 1 . 94E+03 -8.08E-05

SYMBOL SYMBOL CORRESPONDING

PRIORITY PLOTTED STRESS

LOW 5 X STRESS IN X DIRECTION

4 Y STRESS IN Y DIRECTION

3 S SHEAR STRESS IN X-Y PLANE

2 Z STRESS IN Z DIRECTION

HIGH 1 0 ZERO LINE

SIGZ

0.00E-01

4.88E+01

1 . 95E+02

4.39E+02

7.81E+02

1 . 12E+03

1 . 37E+03

1 . 51 E+03

1 . 56E+03

51 E+03

37E+03

12E+03

81E+02

39E+02

95E+02

88E+01

STRESS PLOT: LOAD CASE NO.= 1 MINIMUM STRESS= —0 . 10000E+05

MAXIMUM STRESS= 0.10000E+05

STRESS INCREMENT= 0.43333E+03

THICKNESS

INCREMENT

0 . 25000E-02

I PLY

TYPE

MI N= -0. 1 3000E+05

V

-STRESS- 0 . 13000E+05=MAX

V

+00 .

+90.0

+90.0

+00.0

TTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTT

_l 1 1 (-0 Y 1 1

! 0

-+- -+-

0

0

0Z

0-Z

0

0

0

00

00-Z—Y- 0Z Y

0 Y

0 Y

0 Y- -+- -+-

THICKNESS T ( 1

LISTED DOWN

RIGHT

COLUMN

I

I

I

. 00E-01

. 50E-03

00E-03

. 50E-03

00E-02

25E-02

50E-02

75E-02

00E-02

25E-02

50E-02

75E-02

00E-02

25E-02

50E-02

75E-02

00E-02

COEFFICIENTS OF EXPANSION FOR LAMINATE X-Y COORD

THERMAL COEFF. MOISTURE COEFF

X—D I RECT I ON 1.443E-06 6.147E-04

Y—DIRECT ION 1.443E-06 6.147E-04

SHEAR XY-P LANE 1.422E-12 5.225E-10

! 2

! 5

! 7 -X— ! 1 X ! 1 X ! 1 X I 1 -X— ! 2 X ! 2

! 2

! 2

! 3

! 3

! 3

13

! 4

2 Output data file of results calculated by the laminate computer program ‘ r load case number 1 in Fig. C.l. Stresses are given in tables and graphs.

339
* • * GRAPHITE/EPOXY LAMINATE; LOADS AND STRESSES FOR (0/90)SYM. LAMINATE •••

LOAD CASE NO.- 2 STRESS-FREE TEMP,= 1.00E+02 OPERATING TEMP , =>-3 . 2JE+02 MOIST. CONT.= 0.00E-01

MECHANICALIMID PLANE LAMINATE LOADS

LOAD I NG I NX= 0 00E-01 NY= 0.00E-01 NXY= 0.00E-01

MSW= 1 I MX= 0 00E-01 MY= 0.00E-01 MXY= 0.00E-01

IMIDPLANE LAMINATE STRAINS AND CURVATURES

I EXO= 0. 00E-01 EYO= 0 00E-01 EXYO= 0 00E-01

PLY

ANGLE

I KX= 0.00E-01 KY= 0.00E-01 KXY= 0.00E-01

PLY SIG L SIG T TAU LT SIG X SIG Y TAU XY

10.0

1

-1 . 17E+04 1 . 17E+04 -5.1 IE-04 - 1 . 17E+04 1 . 17E+04 -5.11 E-04

2 90.0 -1 . 17E+04 1 . 17E+04 5. 1 1E-04 1 . 17E+04 - 1 . 17E+04 5.11 E-04

2

3 90.0 -1 . 17E+04 1 . 17E+04 5.11 E-04 1 . 17E+04 - 1 . 17E+04 5.11 E-04

3

40.0 -1 . 17E+04 1 . 17E+04 -5. 1 IE-04 - 1 . 17E+04 1 . 17E+04 -5.11 E-04

4

SIGZ

0 00E-01

294E+02

1 . 17E+03

2 64E+03

4.70E+03

6 . 76E+03

8.22E+03

9. 10E+03

9 40E+03

9 . 10E+03

8 . 22E+03

6. 76E+03

4.70E+03

2.64E+03

1 . 17E+03

2.94E+02

SYMBOL SYMBOL CORRESPONDING

PRIORITY PLOTTED STRESS

LOW 5 X STRESS IN X DIRECTION

4 Y STRESS IN Y DIRECTION

3

S SHEAR STRESS IN X-Y PLANE

2 Z STRESS IN Z DIRECTION

HIGH 1 0 ZERO LINE

STRESS PLOT: LOAD CASE NO.= 2 MINIMUM STRESS= -0 10000E+05

MAXIMUM STRESS= 0.10000E+05

STRESS I NCREMENT= 0.43333E+03

I PLY

TYPE

+00 .

0

+90.0

+90 0

+00.0

THICKNESS T ( I

)

LISTED DOWN

RIGHT

COLUMN

-STRESS > 0 13000E+05=MAX I V I TTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTT I

M I N= -0

V

1 3000E+05

THICKNESS

INCREMENT

0 25000E-02

-+- -+- +- -+ 0- !0Z

! 0

! 0 +0—

! 0

! 0

! 0

-+ 0—

! 0

! 0

! 0

-+ 0—

! 0

! 0

!0Z

-+ 0—

-+-Z-

-+-

Y ! 7

.

-X- ' 1 X ! 1 X ! 1 X ! 1

-Y-I0.00E-01

Y ! 2 . 50E-03

Y ! 5 . 00E-03

. 50E-03

00E-02

. 25E-02

50E-02

, 75E-02

-X-! 2 00E-02

X ! 2 . 25E-02

50E-02

75E-02

00E-02

25E-02

50E-02

75E-02

00E-02

X ! 2 X ! 2

.

-Y- ! 3 Y ! 3

.

Y ! 3 Y

1

3 . -Y- ! 4 .

Fig. C.3 Output data file of results calculated by the laminate computer program for

load case number 2 in Fig. C.l. Stresses are given in tables and graphs.

340
* * * GRAPHITE/EPOXY LAMINATE; LOADS AND STRESSES FOR (0/90)SYM. LAMINATE

LOAD CASE NO.= 3

* • »

STRESS-FREE TEMP.= 0.00E-01 OPERATING TEMP.= 0.00E-01 MOIST. CONT.= 1.00E+00

MECHANICALIMID PLANE LAMINATE LOADS

LOADING I NX= 0.00E-01 NY= 0.00E-01 NXY= 0.00E-01

MSW= 1 I MX= 0.00E-01 MY= 0.00E-01 MXY= 0.00E-01

IMIDPLANE LAMINATE STRAINS AND CURVATURES

I EXO= 0.00E-01 EYO= 0.00E-01 EXYO= 0.00E-01

PLY I KX= 0.00E-01 KY= 0.00E-01 KXY= 0.00E-01

ANGLE

PLY SIG L SIG T TAU LT SIG X SIG Y TAU XY

10,0

1

1 . 02E+04 -1 . 02E+04 4.44E-04 1 . 02E+04 -1 . 02E+04 4 44E-04

2 90.0

2

1 . 02E+04 -1 . 02E+04 -4.44E-04 -1 . 02E+04 1 . 02E+04 -4. 44E-04

3 90.0

3

1 . 02E+04 -1 . 02E+04 -4 . 44E-04 -1 . 02E+04 1 . 02E+04 —4 . 44E-04

40.0

4

1 . 02E+04 -1 . 02E+04 4.44E-04 1 . 02E+04 -1 . 02E+04 4.44E-04

SIGZ

0 . 00E-01

-2.55E+02

-1 . 02E+03

-2.30E+03

-4.08E+03

-5.87E+03

-7. 14E+03

-7.91E+03

-8. 17E+03

-7.91 E+03

-7. 14E+03

-5.87E+03

-4.08E+03

-2.30E+03

-1 . 02E+03

-2.55E+02

SYMBOL SYMBOL CORRESPONDING

PRIORITY PLOTTED STRESS

LOW 5 X STRESS IN X DIRECTION

4 Y STRESS IN Y DIRECTION

3 S SHEAR STRESS IN X-Y PLANE

2 Z STRESS IN Z DIRECTION

HIGH 1 0 ZERO LINE

STRESS PLOT: LOAD CASE NO.= 3 MINIMUM STRESS= -0.10000E+05

MAXIMUM STRESS= 0.10000E+05

STRESS INCREMENT= 0.43333E+03

THICKNESS

INCREMENT

0 25000E-02

I

THICKNESS T ( I ) LISTED DOWN

RIGHT

COLUMN

PLY

TYPE

MI N= -0.

V

1 3000E+05 -STRESS- 0 . 13000E+05=MAX

V TTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTTT

+00.0

+90.0

+90.0

+00.0

-+0- Z0

00

0- 0

00

0- 00

0

0- 0

0 Z0+0

-

-X ——!0 00E-01

X ! 2 50E-03

X ! 5 00E-03

X ! 7 50E-03

-Y — ! 1 00E-02

Y ! 1 25E-02

Y ! 1 50E-02

Y ! 1 75E-02

-Y — !2 00E-02

Y 12 25E-02

Y 12 50E-02

Y ! 2 75E-02

-X — !3 00E-02

X ! 3 25E-02

X 13 50E-02

X 1 3

75E-02

-X — ! 4 00E-02

Fig. C.4 Output data file of results calculated by the laminate computer prr gram

load case number 3 in Fig. C.l. Stresses are given in tables and gra; hr

341



TECHNICAL NOTE

Maurice B. Kasen 1

High Quality Organic Matrix Composite Specimens for

Research Purposes

REFERENCE: Kasen, M. B., “High Quality Organic Matrix Composite Specimens for Research Purposes,” Journal of Composites Technol- ogy & Research, Vol. 8, No. 3, Fall 1986, pp. 103-106.

ABSTRACT: An efficient method for producing and testing organic- matrix composite specimens for research purposes is described. The

production method is adaptable to in-house manufacturing and provides complete control over a large variety of material and processing

variables. The rod-shape, uniaxially reinforced or neat-resin specimens

may be cut to length and tested without further machining. Conventional short-beam shear, flexural strength, and compression test methods may be used. Development of test methods for performing torsional shear tests and for determining the fracture energy G|C are described. Test results at room temperature and at cryogenic temperature are pre- sented. KEYWORDS: cryogenics, fiber-reinforced specimens, fracture energy, neat resin specimens, organic-matrix composites, research specimen

production, short-beam shear, test methodology, torsional shear

Nomenclature

DGEBA Diglycidyi ether of bisphenol A

DDS Diamino diphenyl sulphone

AMD Alkylated methylenedianilene

POPA Polyoxypropylenediamine

FPA Fatty polyamide

BFjMEA Boron trifluoride monoethylamime

degradation of glass reinforced organic matrix electrical insulators. Such insulators are required in the superconducting magnets

of magnetically confined fusion energy systems. Our solution may

interest others who conduct complex composite materials screen- ing programs.

Specimen Selection A small specimen size was dictated by the need to simultaneously irradiate many specimens in a restricted volume and to mini

mize residual radioactivity. The production technique had to he simple and still provide excellent control over the variables of interest. The specimens had to be free of macroscopic defects. It was

desirable to avoid machining costs and the consequent introduction of surface defects. The specimens had to be suitable for test me

in a variety of modes that would provide information on degradation of the fiber-matrix interface. The 3.2-mm (0.125-in.) diameter uniaxially reinforced and neat

resin specimens illustrated on Fig. 1 best met these criteria I . specimen configuration is not new—similar specimens have been

used in studies of reinforced polyester composites However. ^

have adapted the production method to permit its use with. t melt epoxy and bismaleimide matrices and have refined it : duce specimens that are defect free on a macroscopic scale i . 2 illustrates the quality we are able to obtain

Introduction

Experimental screening of the influence of a large number of variables on performance is sometimes necessary when developing

composite materials for a specific service. These variables may in- clude different resin systems and cure states as well as different types and contents of fiber reinforcement. One may wish to study

neat resin performance as well as that of reinforced materials. Unfortunately, the cost of obtaining the large number of well-characterized specimens required for such studies can be prohibitive. We

experienced this problem during an investigation of the combined

influence of cryogenic temperature and neutron irradiation on the

'Staff scientist, Fracture and Deformation Division, National Bureau of Standards, Boulder, CO 80302. © 1986 by the American Society for Testing and Materials

Production Method

Resin-impregnated fiber strands are pulled into a tube n. hi n which they are cured. Although the method is simple in • obtaining specimens free of defects requires careful atten:;< number of details. We find that polytetrafluori*-thv mr I tubes having a very smooth bore produce the best qualm

mens. Glass molds introduce surface defects and rcquirr a remr-e

agent, which might contaminate the specimen surface It is c".m

tial that the glass fiber tows be thoroughly vacuum impregnate . with well-degassed resin. A head of resin must be mamtame

above the fiber during the pulling operation and during 'he • \

2Gauchel, J.. Owens-Coming Fibergla* Technical Center PO H. , ;

' Granville, OH 43023. personal communication

0885-6804/86/0009-0 103*02 50

343
104 JOURNAL OF COMPOSITES TECHNOLOGY & RESEARCH

FIG. 1 —Specimen rods, 3.2 mm diameter, 40 cm long.

MaBttaaaaiaiaa* ...

FIG. 2— Transparency of rods illustrates the quality obtained. Top to bottom: DGBA/DDS. I neatt: DGEBA/DDS. 48 V/O; DGEBA/AMD

/

POPA. (neat): DGEBA/AMD/POPA. 48 V/O; bismaleimide. I neatI.

pressure of about 420 kPa (80 psi) of nitrogen gas must be maintained during the initial stages of the cure cycle.

Proper resin working temperature is maintained by jacketing the impregnation and drawing tubes with a high temperature bath. Neat resin specimens are produced by drawing the degassed resin into the tubes with a vacuum. Reinforced specimens are pulled

from the tube molds, which may be reused. Neat specimens are removed by pressing.

The fiber volume fraction is easily calculated from knowledge of the weight and density of the Fiber and of the tube mold dimensions. The specimen surface is very smooth, and the coefficient of

variability in diameter among rods produced with a given mold is about 0.001 to 0.003. Evolution of solvents in the resin produces porosity in the closed mold. This apparently eliminates many systems of interest where

solvents are conventionally used in the blending process to lower

the melting point of solid cure agents. However, we find that ultra- sonic blending frequently eliminates the need for solvents [ / ]. For example, we routinely blend the solid diamino diphenyl sulphone

(DDS) cure agent into an epoxy resin by this method. Preliminary

experiments indicate that it may be possible to work with solvent- containing resin systems by degassing the impregnated fibers under a vacuum just before pulling into the tube mold. Although

specimens of reasonable quality have been produced in this manner, the technique requires additional development.

Test Methodology

Our current interest is to assess the radiation-induced degradation of the fiber-matrix interface. We have, therefore, concentrated on developing methods for testing in short-beam and tor- sional shear. We have also developed a unique method for measuring the change in longitudinal fracture energy Gic of the specimens. The specimens may also be tested in longitudinal compression using a fixture described elsewhere [2] or in flexure. Each

of these methods requires only that the specimens be cut to the desired length.

The short-beam test method is conventional except for the small specimen size and contouring of the loading nose and supports to match the specimen radius. Specimens are 16 mm (0.6 in.) long for testing at a span/depth ratio of 3.0. Table 1 summarizes apparent

interlaminar shear data obtained by this method at 76 K. The coefficient of variability CV is relatively low among specimens from

different rods of the same material. These data indicate that the short-beam shear test method is relatively insensitive to variation in fiber volume fraction V/O within the range tested. Symmetry of the torsional shear test causes the combined tension, compression, and shear stresses resolved at the fiber-matrix

interface to change more systematically with increasing deformation than is the case for short-beam shear. For a given rotational angle, the magnitude of the shear component is inversely proportional to the gage length and directly proportional to the distance

of the fiber-matrix interface from the specimen center. As shown in Fig. 3a, the failure in this type of test initiates near the specimen

surface and progresses inward. Figure 3b illustrates the interfacial nature of the failure mode. Table 2 summarizes the ultimate tor- sional strengths and coefficients of variation obtained during test- ing three epoxy matrix systems at 7b K. The torsional specimen is 89 mm (3.5-in.) long. 35 mm (1.375

in.) of which is embedded in 9.52-mm (0.375-in.) diameter end

caps, leaving a gage length of 19 mm (0.75 in.). Specimens are ep- oxy bonded into linen phenolic end caps for room temperature test- ing. For cryogenic testing, we freeze the specimens into aluminum

end caps to permit recovery of the specimen ends for additional

testing. Several types of aqueous slurries provided sufficient bond

strength to fracture the specimen. However, progressive cracking

of the bond from the gage section into the end caps increased the

effective gage length, complicating calculation of the torsional modulus. This was corrected by incorporating a senes of four set screws into the end caps as illustrated on Fig. 4. Contact of these

TABLE 1 —Short-beam shear data. 76 K. DGEBA

FPA. Type E Glass (10 specimens each condition!.

V/ O Glass Rod

Strength.

MPa CV

41 1 23.6 0.042

2 23.0 0.087

23.3 0.064

48 i 23.3 0.047

2 23.6 0.013

23.5 0.034

344
KASEN ON ORGANIC MATRIX COMPOSITE SPECIMENS 105

FIG. 3— Cross sections of 48 V/O glass reinforced specimen after tor- sion testing at 76 K. (a) Full cross section illustrating initiation offracture

at the surface and propagation inward, (b) Magnified view of the fracture showing the interfacial nature of the crack propagation. Original magnification X 1000. FIG. 4—Gripping system for cryogenic torsional shear test.

TABLE 2— Torsional shear data. 76 K. (5 specimens each condition, 48 V/O Type E glass).

Epoxy Matrix Strength, MPa CV

DGEBA/DDS/BFjMEA 158.4 0.086 DGEBA/DDS 160.8 0.036 DGEBA/AMD/POPA 148.6 0.078

screws with the specimen prevents rotation and confines the failure

to the gage section. The torsional load is applied by a stainless steel shaft that mates with a slot in the upper end cap, while the lower end cap is pinned to a stainless steel reaction tube. Although the above tests are useful in a materials screening pro- gram, they provide only qualitative information on the integrity of the fiber-matrix interface. We are, therefore, developing a test method that provides information on the longitudinal fracture en- ergy Gic

. The principle is illustrated on Fig. 5a. A razor blade is pressed an arbitrary distance into the end of the fiber reinforced specimen causing a crack to propagate lengthwise down the speci- men. The fracture energy is then calculated by the equation

GIc = 0.01372 d

i Ey2/x* (1)

where E is the Young’s modulus and where d, y and x are as de- fined on Fig. 5a. This approach is based on the original work of Obreimoff [ J] who developed a similar expression for the splitting

strength of mica and on the subsequent work of Barlow and Windle [4], who suggested that the basic method could be applied to uniaxially reinforced composite materials. However, Eq 1 differs from that of Obreimoff in that the moment of inertia has been cal- culated about the centroidal axis of a semicircular cross section to take into consideration the geometry of the rod-shape specimens.

Obreimoff defined the fracture energy Q as the energy release re- quired to create one of the two new surfaces. This was equated to G|C by Barlow and Windle, and we have followed this convention

rather than that in which Gic is taken as the energy required to create both surfaces. The prototype fixture illustrated in Fig. 5 b was used to generate

the G|C values presented in Table 3 for several materials and temperatures. The results are in reasonable agreement with the 1.9 to

FIG. 5— Test method for determining fracture mtrgf G

diameter fiber-reinforced specimens, (a) Schematic "t the .]•••., .< ing the terms in Eq I. Blade thickness 2y was 0.035 mm

present work, (b) Prototype fixture used for fracture <-nrrg\ ttttt plunger restricts blade penetration to 4 mm

TABLE 3 —Fracture strength G i, 4H V O Type F giu\s lone specimen each).

Resin System T. K z. mm G„. kJ m‘*

DGEBA/FPA 76 6.0 2.09 DGEBA/DDS/BFjMEA 76 9.0 0.41 DGEBA/AMD/POPA 76 8.5 0.52 DGEBA/AMD/POPA 295 6.5 1.52

345
106 JOURNAL OF COMPOSITES TECHNOLOGY & RESEARCH

3.2 kJ • m

-2 values reported for the fracture energy of similar materials, as determined by the double cantilever beam method [4], suggesting that this may be a useful test. The direction and magnitude of the changes are also reasonable for the differing materials and test temperatures.

International Cooperation

Development of the described production method and test methodologies is currently being pursued by NBS, the Institute for

Scientific and Industrial Research, Osaka University, Japan, and

the Rutherford Appleton Laboratories in England, The Japanese

group has successfully applied the technique to specimens rein- forced with graphite, silicon carbide, and aluminum oxide fibers in addition to glass [5, 6], The objective is to achieve a commonality in specimen configuration and testing procedure that will increase the effectiveness of technology transfer between laboratories having common research goals.

Acknowledgment

This work was sponsored by the Office of Fusion Energy, U.S. Department of Energy. The author acknowledges the assistance of

P. T. Purtscher and R. Stoddard in specimen preparation and test- ing. The contribution of R. D. Kriz in deriving Eq 1 is particularly

appreciated.

References

[7] Hodges, W. T. and St. Clair, T. L., "Ultrasonic Mixing of Epoxy Curing Agents,” Technical Memorandum 85643, NASA Langley Research

Center, Hampton, VA, 1983.

[2] Schramm, R. E. and Kasen, M. B., "Cryogenic Mechanical Properties

of Boron-, Graphite-, and Glass-reinforced Composites", Materials Science and Engineering , Vol. 30, 1977, pp. 197-204.

[J] -OJjreimoff, J. W., “The Splitting Strength of Mica", Proceedings of

the Royal Society, Vol. A127, 1930. pp. 290-297.

[4] Barlow, C. Y. and Windle, A. H., “Razor Blade Test for Composite

Toughness," Journal of Materials Science and Letters, Vol. 4. 1985, pp. 233-234.

[5] Okada, T., Nishijima. S., Yamaoka, H.. Mivata. K.. Tsuchida. Y.. Mizobuchi, K., Kuraoka, Y.. and Namba, S., "Mechanical Properties

of Unidirectionally Reinforced Materials." Sonmetallic Materials and

Composites at Low Temperatures—3, G. Hartwig and D. Evans. Eds.. Plenum Press. New York, in press.

[6] Takeno, M., Nishijima. S., Okada. T.. Fujioka. K.. Tsuchida. Y.. and Kuraoka, Y., "Thermal and Mechanical Properties of Advanced

Composite Materials," Advances in Cryogenic Engineering t Materials), Vol. 32, Plenum Press, New York. 1986. pp. 217-224.

346

II
STRATEGY FOR THE DATA BASE CONSTRUCTION ON RADIATION -RESISTANT

CRYOGENIC COMPOSITE INSULATORS FOR MAGNETIC FUSION ENERGY APPLICATIONS*

Maurice B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

A strategy is suggested for the development of fiberreinforced, organic insulators to be used in superconducting

magnets for magnetic fusion energy power systems. Two data

bases are required. The first is a component data base providing information for basic material selection. The second is an engineering data base generated on insulators fabricated from

the selected components. Successful completion of these tasks

requires multidisciplinary expertise. Cooperative research is

presently under way among laboratories in the United States,

Japan, and England.

INTRODUCTION

The conversion of fusion energy to electrical power by magnetic con- finement of the plasma will require the use of superconducting magnets.

Such magnets must function without repair over the lifetime of the installation. The large magnet size will create substantial electromagnetically

induced forces on the structural components. Furthermore, portions of the

magnets will have to sustain an accumulated irradiation dose equivalent to

100 to 1000 MGy (MGy = 10 8 rads), if excessive shielding costs are to be

avoided. 1 More than half the dose will come from neutrons.

Under these conditions, the performance of a number of magnet com- ponents, such as the superconducting elements and stabilizing materials,

is uncertain. However, it appears likely that the permissible fluence wi 1

be primarily limited by radiation- induced degradation of the polymeric materials that conventionally serve as electrical and thermal insulators. " The

use of inorganic insulators could alleviate this problem but they might be

prohibitively expensive.

*Published in Proceedings, International Symposium on Fundamental Rcscar:,

Strategy in the Development of New Materials for Efficient Energy Conversion. Osaka, Japan: University of Osaka; 1986. pp . 112-118.

347
Therefore, it is necessary to define the radiation tolerance and potential insulation capabilities of organic materials and to develop the most

promising materials. These organic insulators are unreinforced polymer

films and sheets, potting compounds, and fiber-reinforced composite products. The fiber-reinforced insulators are particularly critical because

they also serve as structural elements.

This document therefore addresses the strategy for efficient development of the fiber-reinforced insulating materials. Although the specific

ideas expressed are those of the author, they reflect many discussions held

with other individuals, especially with Prof. T. Okada and Dr. S. Nishijima,

Institute of Industrial and Scientific Research (ISIR), Osaka University,

Japan, and with Dr. David Evans, Rutherford Appleton Laboratory, England.

APPROACH TO THE PROBLEM

Successful development of fiber-reinforced composite materials for such

demanding service requires knowledge of the intrinsic performance of the

composite constituents under the prevailing environmental conditions. It

is the basis for optimal selection of the individual components from which

functional insulating materials are to be fabricated. It is also required

for determining the parameters that must be controlled during fabrication

to achieve minimal variability in performance. The characterization of

practical insulating laminates produced with this component data base will

provide the engineering data base required by magnet designers.

This effort requires the cooperation of individuals having expertise

in polymer science and physics, composite material fabrication, cryogenic

testing, and cryogenic radiation science. The fundamental strategy envisions close cooperation among laboratories having these areas of expertise.

Laboratories in the United States (NBS, ORNL), in Japan (ISIR, Osaka University), and in England (Rutherford Appleton Laboratory) are presently

cooperating in this effort.

THE COMPONENT DATA BASE

Establishing a component data base requires an extensive materials

screening program, which must be carefully planned to avoid unnecessary

effort and to minimize cost. A significant element in the strategy is to

ensure maximum efficiency in technology transfer among laboratories co- operating in such research. The first step, a common approach to specimen

design, has been completed.

Soecimen Standardization —A

The many variables that need to be evaluated mandate that the studies

be conducted with small, highly characterized, unreinforced resin and fiberreinforced specimens that are easily produced and tested. Their fabrication

process must be simple, inexpensive, and very flexible so that the parameters to be studied can be systematically altered. Machining of specimens,

348
which is expensive and introduces surface flaws, must be avoided. The

configuration must not be susceptible to edge effects, which could invalidate test data; at the same time, the specimens must be suitable for testing

in a variety of modes.

These considerations have led to international agreement that the

optimum specimen configuration would be a rod approximately 3.2 mm in

diameter, produced by a variation of the vacuum pultrusion process. Techniques for production of such specimens have been successfully developed in

Japan, 3 England, 4 and the United States. 5 Typical specimens are shown in

Fig. 1.

Mechanical Testing

Efficient test methods for use with this standard specimen design are under cooperative development. Emphasis is on mechanical property testing

because it is commonly accepted that the dielectric strength of insulators

is degraded only after the mechanical performance has been degraded. 6 Therefore, electrical insulating performance is not included in the com- ponent data base but it will be a part of the engineering data base.*

The specimens, cut to appropriate lengths, will be tested by a variety

of methods. Flexural tests are traditional and easily performed. However,

lack of symmetry introduces a complex and changing stress state throughout

the test, making it difficult to quantify the results. The high failure

strain of glass-reinforced specimens at cryogenic temperatures further

complicates interpretation of the test data. 3 ' 7 The short-beam shear test

appears to be useful for rough screening of the integrity of the fibermatrix interface, and this test method requires very little specimen

material. 8 However, the data remain qualitative.

A cryogenic torsional shear test method developed at NBS may prove to

be very useful in this program. 8,9 Figure 2(a) illustrates a typical

torsional stress-strain curve obtained at 4 K on a glass-fiber-reinforced

specimen. The test method generates quantitative information on the modulus, yield strength, and ultimate strength. Progressive damage accumulation

is evidenced by successive load drops. Figure 2(b) illustrates the interfacial nature of the failure path. This is advantageous because available

data suggest that degradation of this interface may be a dominant factor in

overall mechanical property degradation during cryogenic irradiation. 10 The

torsional test method has also been applied to unreinforced resin specimens

at 4 K.

A test procedure for measuring the longitudinal fracture energy, T. , of the fiber-reinforced rod specimens at cryogenic temperatures has also

been developed at NBS. 5 As illustrated by Fig. 3(a), the method is very

*A three year research program on the dielectric strength of insulators at cryogenic temperatures will start at ISIR in 1986 as a program comph'rv’ntary to that discussed in this paper.

349
simple, requiring only measurement of the length of the crack that is

produced by pressing a razor blade into the specimen end. The sensitivity

of this method is illustrated in Fig. 3(b). With this test procedure,

changes in the integrity of the fiber-matrix interface due to irradiation

can be associated directly with a change in the fracture energy.

To obtain the maximum amount of information from a minimum number of

specimens, NBS has developed a method for recovering the undamaged portions

of the torsional test specimen after the test has been completed . 9 Each

specimen end may then be subjected to a fracture strength test plus a shortbeam shear test.

Failure Mode Characterization

Fractographic analysis by scanning electron microscopy contributes to an understanding of how radiation influences the failure mode. However, a full understanding of the radiation degradation process requires correlating

the degradation in mechanical properties with specific damage induced on a molecular level. This will require cooperation of individuals skilled in

the use of advanced analytical techniques in polymer chemistry and physics.

The results will contribute much to basic polymer science and will provide a firm scientific basis for selecting polymer systems for additional study.

Component Parameters

Selection of matrix resin systems for this study should be made by

individuals with a background in polymer chemistry and familiarity with the

current state of knowledge about the influence of ionizing irradiation on polymer degradation. Several recent publications have addressed parameters

that must be considered . 11-13

The influence of resin chemistry must be carefully evaluated. Catalysts added to reduce processing time frequently contain boron compounds.

It must be determined whether the resulting increase in neutron capture

cross section contributes significantly to the level of degradation. The

significance of trace metallic elements on resin performance should also be

evaluated, and the influence of the resin cure state on resin degradation

should be determined.

The type E glass fiber normally used in insulator construction contains

10 to 12% B 2 O 3

. It must be determined whether the higher neutron capture

cross section of this type of fiber is sufficiently detrimental to warrant

the additional expense of fabricating insulators with boron- free glass

fiber. The radiation performance of specimens reinforced with new fibers,

such as AI 2 O 3 and SiC, should also be evaluated. This work has already

begun at ISIR . 1

4

Dependence of radiation resistance on fiber volume fraction must be

studied to distinguish between the amount of damage sustained by the resin

and by the resin-matrix interface. This information is needed to establish

the optimum fiber content for maximum radiation resistance.

350
The high hydrogen content in the surface finishes normally applied to

glass fibers to improve the fiber-matrix bond strength is expected to con- centrate neutron- induced damage at the fiber interface. 1 This must be

evaluated to determine whether control of the surface finish chemistry is

necessary in fabricating functional insulating laminates.

Radiation Parameters

Selection of the irradiation conditions must be made by individuals

having expertise in radiation science and knowledgeable about the anticipated environment in a fusion reactor. There must be access to a reactor

capable of providing the required fluence and spectra at 4 K. The required

radiation environment can be simulated to some extent in a fission reactor

where neutron and gamma rays coexist. However, few fission reactors have

facilities for irradiation at liquid helium temperature, and the available

fluence is inadequate to provide the required dosage in a reasonable time.

Completion of the National Low Temperature Neutron Irradiation Facility

(NLTNIF) facility at the Oak Ridge National Laboratory (ORNL) will provide

the necessary irradiation environment for this work. 16 Scheduled to become

operative in the summer of 1986, this facility will provide the required

fluence level of about 10 18 neutrons/cm2 in a relatively short time as well

as providing a relatively large irradiation volume.

The exact radiation spectra that will exist at the superconducting

magnets in a functioning MFE system are not known. Therefore, it will

be valuable to assess damage accumulation and property degradation as a function of a known change in radiation spectra. This can also be accom- plished in the NLTNIF reactor.

Nevertheless, radiation research at 4 K is expensive and requires

unique facilities. Therefore, the extent to which studies conducted at

higher temperatures are relevant to 4-K performance should be determined.

Damage accumulated during irradiation at 4 K must be distinguished from

that occurring due to the release of stored energy during warmup. This is

important because the magnets will be periodically warmed up during their

lifetime. In situ damage assessment will require development of techniques

for testing either in the NLTNIF cryostat irradiation zone or in the test

chamber located immediately above the irradiation zone. Alternatively,

irradiated specimens could be transferred at 4 K from the NLTNIF facility

to a remote test where they would then be transferred without warmup into

the test apparatus. Neither task will be easily accomplished.

CURRENT PROGRAM STATUS

One hundred sixty-five specimens provided by NBS will be Irradiated

in the NLTNIF facility as soon as it becomes operative. These consist of

unreinforced resin specimens and specimens reinforced with 48 volume pen . : of type E glass. One set of specimens duplicates the commercial G-11CR

formulation to provide a baseline of performance against which the performance of the other materials can be compared. Another series of sperim<-:.-.

351
duplicates the G-11CR formulation except for the BF3 MEA accelerator. A

third set was fabricated with an epoxy matrix typical of that which would

be used for impregnation or potting purposes. A fourth set of specimens

was fabricated with a state-of-the-art bismaleimide resin.

The specimens will be irradiated at ORNL to three increasing fluence

levels at 4 K. The specimens will be warmed to room temperature and re- turned to NBi’ for 4-K testing. A duplicate set of unirradiated materials

will also be tested at 4 K.

Following this initial test run, a second irradiation program will be

undertaken with specimens provided by ISIR. NBS will submit these specimens

to ORNL for irradiation, after which they will be returned to NBS for testing by visiting personnel from ISIR.

SUMMARY

The goal of developing functional organic insulators required by superconducting magnets in MFE reactors requires a systematic, interdisciplinary

research program following a well-defined strategy. The initial objective

is to develop a data base defining the optimum components for the required

service. This data base can then be used in the design and fabrication of

engineering laminates required for specific applications in magnet construction. The characterization of such laminates will provide an engineering

data base for the designer. This document presents the elements of the

strategy necessary to attain these goals.

The magnitude of the task at hand and recognition of the various

disciplines required for its successful accomplishment has stimulated

cooperation between laboratories in the United States, Japan, and England.

Significant progress has already been made in standardizing test specimens

and test methodology. Cooperation will soon extend to the irradiation

programs

.

ACKNOWLEDGMENT

The work at NBS is sponsored by the Division of Magnetic Fusion Energy,

U.S. Department of Energy, Washington, D.C.

REFERENCES

1. G. F. Hurley and R. R. Coltman, Organic Materials for Fusion Reactor

Applications, J. Nucl. Mater. 123 and 124: 1327-1337, 1984.

2. B. S. Brown, Radiation Effects in Superconducting Fusion-Magnet Materials, J. Nucl. Mater., 97: 1-14, 1981.

3. T. Okada, S. Nishijima, H. Yamaoka, K. Miyata, Y. Tsuchida,

K. Mizobuchi, Y. Kuraoka, and S. Namba, Mechanical Properties of

Unidirectionally Reinforced Materials, in "Nonmetallic Materials and

Composites at Low Temperatures -- 3

, " G. Hartwig and D. Evans, eds

,

Plenum Press, New York (in press).

352
4. D. Evans, Rutherford Appleton Laboratories, Chilton, Didcot, Oxon,

England (private communication, 1983)

5. M. B. Kasen, High Quality Organic Matrix Composite Specimens for

Research Purposes, J. Compos. Technol. Res. (accepted for publication).

6. R. H. Kernohan, C. J. Long, and R. R. Coltman, Jr., Cryogenic Radiation

Effects on Electrical Insulators, J. Nucl. Mater. 85 and 86: 379-383,

1979.

7. T. Okada, S. Nishijima, and H. Yamaoka, Radiation Damage of Composite

Materials- -Methods and Evaluation, in "Advances in Cryogenic Engineering- -Materials , " 32, R. P. Reed and A. F. Clark, eds . , Plenum Press,

New York, 1986, 145-151.

8. M. B. Kasen and R. B. Stoddard, Screening the Performance of Organic

Insulators under Cryogenic Neutron Irradiation, in "Advances in Cryogenic Engineering- -Materials , " 32, R. P. Reed and A. F. Clark, eds.,

Plenum Press, New York, 1986, 153-159.

9. M. B. Kasen, Cryogenic Torsional Mechanical Test Apparatus, submitted

for publication to Cryogenics.

10. H. W. Weber, Low Temperature Neutron and Gamma Irradiation of Glass

Fiber Reinforced Epoxies, J. Nucl. Mater., 115: 11-15, 1983.

11. D. Evans and J. T. Morgan, A Review of the Effects of Ionizing Radiation on Plastic Materials at Low Temperatures, in "Advances in Cryogenic Engineering- -Materials , " 28, R. P. Reed and A. F. Clark, eds.,

Plenum Press, New York, 1982, 147-164.

12. H. M. Banford, Electrical Insulation and Fusion Reactors, in "Advances

in Nuclear Science and Technology," 16: 1-73, J. Lewins and M. Becker,

eds., Plenum Press, New York, 1984, 1-73.

13. W. Maurer, Neutron and Gamma Irradiation Effects on Organic Insulating

Materials for Fusion Magnets, Report KFK 3974, Institut fur Technische

Physik, Kernforschungszentrum Karlsruhe, Germany, 1985.

14. M. Takeno, S. Nishijima, T. Okada, K. Fujioka, Y. Tsuchida, and

Y. Kuraoka, Thermal and Mechanical Properties of Advanced Composite’

Materials at Low Temperature, in "Advances in Cryogenic Engineer i ng

Materials," 32, R. P. Reed and A. F. Clark, eds., Plenum Press, New

York, 1986, 217-224.

15. R. R. Coltman, Jr., C. E. Klabunde, and F. W. Young, Jr.

A National

Low Temperature Neutron Irradiation Facility, J. Nucl. Mater., 122 a: 123: 1357-1358, 1984.

353
Fig. 1. Specimen rods, 3.2-mm diameter and 40-cm long,

containing 48 volume percent type E glass in an epoxy matrix. Similar specimens are produced from

unreinforced epoxy resin.

334
(b)

Fig. 2. (a) Typical torsional stress-strain curves for a glass

fiber-reinforced specimen obtained at 4 , 77, and 295 K. (b) Cross section of the failed specimen illustrating

the interfacial nature of the crack path.

355
Fig. 3.

Blade

(b)

(a) Method for determining the fracture energy, Gj c , on 3. 2 -mm diameter, fiber-reinforced specimens. The equation was derived

for the specific specimen configuration. 5 E - Young's modulus.

(b) Difference in crack length observed at 295 K (upper specimen)

and at 77 K (lower specimen). Their Gj c values are 1.52 and

0.52 kJ-m"2, respectively.

356
TECHNOLOGY TRANSFER

357

TECHNOLOGY TRANSEER PROGRAM

LEADER: N. J. Simon

CONTRIBUTORS: L. Delgado, C. J. King, R. P. Reed

OBJECTIVES

• Organization of workshops to promote interaction between designers

and material specialists, to discuss issues related to low- temperature

material needs of the fusion energy program, and to present new lowtemperature data for structural alloys, composites, and weldments

• Preparation of an annual report, as well as monthly highlight reports

to the Office of Fusion Energy, U.S. Department of Energy

• Evaluation of low- temperature mechanical and physical properties of con- ductor and structural materials for cryogenic copper and superconducting

magnets; preparation of handbook pages and supporting documentation; distribution of handbook pages to participants in the fusion energy program

and to the "Materials Handbook for Fusion Energy Systems"

RESEARCH HIGHLIGHTS

• The Ninth Annual Cryogenic Structural Materials Workshop was organized;

it was held October 6—8, 1986 in Reno, Nevada in conjunction with the

United States—Japan Low Temperature Structural Materials and Standards

Workshop on October 9—10.

• "Materials Studies for Magnetic Fusion Energy Applications-IX," (NBSIR

86-3050, 338 pages, 1986, editor R. P. Reed) was prepared, published,

and distributed.

• Data were collected for handbook pages on C10100—C10700 copper and

C17000—C17510 beryllium copper. Over one thousand documents (reports

and journal articles) were acquired, coded for property information,

and entered into a data base management system. Evaluation of document

pertaining to C10100—C10700 tensile, electromagnetic, creep, and f at

properties has been completed.

• Handbook pages and supporting documentation covering tensile and electromagnetic properties of C10100-C10700 copper have been distributed.

Cryogenic C10100—C10700 data on creep and fatigue were evaluated and

reviewed together with tensile and electrical property data at the

Ninth Annual Cryogenic Structural Materials Workshop.

359

NINTH ANNUAL

CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

Airport Plaza Hotel

Reno , Nevada

October 6—8, 1986

This annual workshop provides an opportunity for discussion

of problems, advances, and goals in the development of lowtemperature materials for cryogenic and superconducting magnets

for fusion energy systems. The 1986 workshop emphasized:

• Cryogenic properties of copper alloys and laminates

pertaining to design requirements for the compact

ignition tokamak (CIT)

• Identification of materials needs for future magnet

systems

• New low- temperature materials developments in the

United States and Japan and international research

cooperation

361

9th annual

Cryogenic Structural Materials Workshop — Program

MONDAY

OCTOBER 6

8:30 a .m. Introductions R. P. Reed,

and V. Der,

NBS , OFE

COMPACT IGNITION TOKAMAKS

8:45 a . m. Toroidal Field Coil Design J. Citrolo, PPPL

9:30 a .mDesign of Alcator C-Mod H. Becker, MIT

10:00 a .mu Copper Alloy and Composite Properties

Review N . J . S imon

NBS

Mechanical Properties of Copper L. L. Summers, LLNL

Mechanical and Physical Properties

of Copper Alloys

R. P. Reed, NBS

Composite Properties H. Becker, MIT

Laminate Data L. L. Scull, NBS

12:15 p .m. Lunch

1:30 p .m

Discussion of Structural Alloy Research

Needs for Compact Ignition Tokamaks

NONMETALLIC STRUCTURAL COMPOSITES

2:30 p ,m. Edge Stresses in Woven Laminates R. D. Kriz, NBS

3:00 p .mCharacterization of Cryogenic Composite

Materials

M. K. Abdelsalam,

U. Wisconsin

3:30 p .mCryogenic Nonmetallic Composite Research

at ISIR

S. Nishijima,

Osaka U.

4:00 p .m. Composite Test Method Development M. B. Kasen, NBS

4:30 p • m • Low Temperature Irradiation Facility H. R. Kerchner, ORN

TUESDAY, OCTOBER 7

LARGE-SCALE TOKAMAK AND MIRROR FUSION ENERGY DEVICES:

DESIGN AND MATERIALS

9:00 sl • m • Advanced Magnet Designs and Materials

Problems

J. R. Miller, LLNL

9:45 a.m. NET Plans and Related Materials Research A. Nyilas, KFK

10:30 a .m. LCP Experiments T. McManus, ORNL

11:00 a • m • Material Selections for the SSC T. H. Nicol,

Fermi lab

11:30 a .m. Activation of Structural Alloys

Fusion Reactor Magnets

in F. M. Mann,

Westinghouse Hanfor

363
TUESDAY, OCTOBER 7, continued

12:00 m. Lunch

2:00 P « III e Modification of 316LN-Type Alloys P. T. Purtscher, NBS

2:30 p .m c Design of 316LN-Type Alloys N. Jc Simon, NBS

3:00 p .m. Sources of Intergranular Embrittlement

in Fe-Mn Alloys

M. Strum, LBL

3:30 p .m. Low-Coef ficient-of- Expans ion

Structural Materials Development

for Conductor Applications

R. Ballinger, MIT

4:00 p o m • Preliminary Test Results of Magnetic

Field Effects on Fracture Toughness

H. Takahashi and

T. Shoji, Tohoku U.

4:30 p «m. Mechanical Properties of Al-Li 2090-T8 A. S unwoo, LBL

6:00 p 0 m. Reception

WEDNESDAY, OCTOBER 8

LARGE-SCALE TOKAMAK AND MIRROR FUSION ENERGY DEVICES:

MATERIALS AND WELDING

8:30 a.m.

9:00 a.m.

9:30 a.m.

10:00 a.m. 10:30 a.m.

11:00 a.m.

11:30 a.m.

12:00 m.

Effects of Strain Rate on Mechanical

Properties of Austenitic Stainless

Steels at 4 K

The Effect of Inclusions and Nickel on Weld Toughness and Fracture Appearance

High Compliance Fracture Toughness Test

at 4 K for 310S Stainless Steel and

Its Weldments

Welding of Austenitic Manganese Steels

Manganese, Molybdenum, and Nitrogen in

Stainless Steel SMA Welds

Mechanical Properties of Welded Joints

of Newly Developed Stainless Steel

Extension of Mechanical Property Test

Standards of Metals and Composites to

Liquid Helium Temperature

Discussion of Structural Alloy Research

Needs for Large Superconducting Magnets

Y. Kohsaka, NKK

T. A. Siewert, NBS

H. Takahashi,

Tohoku U.

J. W. Chan, LBL

C. N. McCowan, NBS

K. Yoshida, JAERI

B . M. S trauss , Teledyne Eng.

THURSDAY and FRIDAY, OCTOBER 9 and 10

UNITED STATES-JAPAN WORKSHOP

on Low Temperature Structural Materials and Standards

364
NINTH ANNUAL CRYOGENIC STRUCTURAL MATERIALS WORKSHOP

SUMMARY OF TECHNICAL PRESENTATIONS

SESSION I. COMPACT IGNITION TOKAMAKS

Because of the planned construction of an intermediate ignition fusion

reactor with copper or copper alloy coils operating at 77 K, the workshop

this year opened with a session on design requirements and properties of

copper and copper alloys. J. Citrolo (PPPL) reviewed the design features

of the Compact Ignition Tokamak (CIT) toroidal field coils that will attain

a field of 16 T. Since operating stresses in the copper are expected to be

-490 MPa (70 ksi) , considerably above the yield strength of the cold-worked

C10700 copper conductor, a current design places the conductor between

Inconel 718 plates, to which it will be explosively bonded. Design requirements of another compact copper-coil device, the Alcator C-Mod, were discussed by H. Becker ( MIT). This machine will be held together by vertical

tie rods (nonwelded) , between 316L and 316LN cover plates; it will experience 50,000 pulse cycles over its expected period of operation. Therefore,

the device is being designed to meet fatigue criteria, and provisions for

periodic inspection will be included in the design. (The CIT is designed

for only 3000 pulses.) Following a discussion of design requirements for

these copper machines, a review of the C10100—C10700 copper data base was presented by N.J. Simon of NBS. One important conclusion of the review is

that long-term creep data at room temperature and below are not available.

L. L. Summers (LLNL) reviewed results of a copper alloy testing program in

which the C10400, C15500, and C17510 alloys were evaluated. A significant

result is that at room temperature, the magnitude of the C15500 ultimate

tensile strength approaches that of the yield strength, which results in

less safety margin than is usually provided when the design allowable

stresses are set as a fraction of yield strength. Results of extensive

testing of C10400, C15500, and C17510 in the annealed, half -hard, and hard

conditions were presented by R. P. Reed (NBS). Both yield and tensile

strength were found to increase significantly below room temperature for

material in the hard condition. Tensile test results, measurements of the

electrical resistivity as a function of temperature, and the time -dependent:

properties (creep, fatigue, and fatigue—creep) at 77 K and ambient temperature were presented. The session continued with reports by H. Becker (MIT

and L. Scull (NBS) on the testing of laminated copper. Under compression

loading, copper extruded from the Inconel—C10200 laminates tested at MIT

For the roll-bonded copper—stainless steel (Fe—18Cr—13Mn—3Ni) laminates

tested at NBS, reasonable agreement was found between the measurements of

yield strength and Young's modulus and predictions based on mixture rules

and copper and stainless steel properties.

A discussion session on materials requirements for compact ignition

tokamaks followed. Reasons for the selection of a laminate re inforr eyr.er.- to compensate for the inadequate yield strength of half -hard C10700 oopp«-r

were discussed. Since other high-conductivity copper alloys approach • :.<•

365
required strengths, other considerations, such as thermal conductivity,

entered into the selection process in which a complex computer code was utilized. Further discussion indicated some uncertainty over the amount

of cold work that 19 -mm (3/4- in) thick C10700 plate could be subjected to.

The planned test matrix for the laminate material was outlined by G. Brown

(PPPL) . Since the matrix is extensive, and therefore it will be expensive,

comments were solicited from representatives of the various national

laboratories

.

SESSION II. NONMETALLIC STRUCTURAL COMPOSITES

New results from modeling studies of woven glass—epoxy laminates such

as G-10CR and G-11CR were presented by R. D. Kriz (NBS) . Finite-element

results demonstrated that the weave geometry reduces edge stresses at low

temperatures, in accord with previous experimental results that showed

that most of the damage occurs in the interior of these laminates. M. K.

Abdelsalam (University of Wisconsin) presented an overview of property

measurements on several different load-bearing thermal insulators to be

used in compression. Thermal conductivity, fracture, and fatigue studies

were discussed. The extensive facilities of the Institute of Scientific

Industrial Research (ISIR) at Osaka University, Japan, were described by

S. Nishijima. Their facilities are capable of conducting flexural, tensile,

compressive, impact, creep, thermal, electrical, and radiation tests at

cryogenic temperatures. An adhesive bonding method for producing composite

cryostats of complicated shape has been developed. M. B. Kasen (NBS) described the development of well-characterized small specimens for a screening program that will include the effects of low- temperature irradiation.

The National Low-Temperature Neutron Irradiation Facility (NLTNIF) in which

these specimens will be irradiated was described by H. R. Kerchner (ORNL)

.

Specimens can be neutron irradiated at 3.2 K in a 12-T field and tested

without warmup.

SESSION III. IARGE- SCALE TOKAMAK AND MIRROR FUSION ENERGY DEVICES:

DESIGN AND MATERIALS

An overview of the superconducting magnet development program at LLNL

was presented by L. L. Summers. Expected future use of an advanced superconductor, NbN, for which J c is much less sensitive to degradation from

strain or irradiation than that of Nb 3 Sn and NbTi superconductors in current

use, is now driving some of the materials selection processes. The TIBER II

steady-state ignition tokamak, a very compact machine with a high radiation

fluence and very high stress levels, also places increased requirements on materials. However, 316LN may fulfill magnet case requirements. The effects of thermomechanical processing schedules on conduit alloys for a superconducting cable- in-conduit magnet configuration were also discussed.

Plans for the NET (Next European Torus) were presented by A. Nyilas

(KFK) . This machine will be intermediate between the existing JET (Joint

European Torus) and the larger DEMO. NET will burn deuterium and so provide

366
materials testing in a reactor environment and also a test of tritium handling capabilities. Laser welding will be used for the superconducting coil

conduit in NET. A program was described for scaling up tests on small specimens so that stresses in complicated geometries can be identified. Again,

316LN is the chief structural material under consideration, which reinforces

the relevance of recent NBS fundamental studies of this alloy.

T. McManus (ORNL) described recent testing of the General Dynamics,

General Electric, and Westinghouse superconducting magnets for the Large

Coil Project (LCP) , a collaborative effort between the United States,

Europe, and Japan. Six separate D-coils have been provided for testing.

The GD and GE coils use helium pool boiling: both showed degraded heat

transfer during quench testing, which is believed to be due to vapor accu- mulation. The Westinghouse coil showed a significant reduction of critical

current that was assumed to be due to straining of the Nb 3 Sn conductor

during fabrication. In spite of these imperfections, the overall results

of the LCP are encouraging, since all magnets were successfully tested as single coils with peak fields near 6 T.

The focus changed with a presentation by T. H. Nicol (Fermilab) on materials selection considerations for the Superconducting Super Collider

(SSC) . This facility, although not a fusion device, will have about ten

thousand superconducting magnets in an underground tunnel about one hundred

kilometers in diameter. Design considerations are based on the successful

Tevatron built by Fermilab, another large-scale device with more than one thousand superconducting magnets and similar refrigeration requirements.

To keep operational costs low, the SSC will have a relatively small refrigeration load for a device of its size. A key feature of the SSC is a support post of reinforced composite designed to minimize the heat leak.

Reliable information on low- temperature reinforced composite properties is,

therefore, an important consideration.

F. W. Mann (HEDL) discussed the use of the existing Starfire conceptual

design code to determine neutron flux and to predict transmutation activation for a series of representative alloy compositions supplied by NBS. The

neutron flux spectrum at the magnet was found to be relatively soft; thus,

the conclusion was that there should be no significant activation problems

regarding short-term maintenance or long-term disposal requirements. However, T. Okada (Osaka University) reached the opposite conclusion with tin- assumption of 10% hard (14 MeV) neutrons in the spectra. Thus, the activation problems would be similar to the first-wall conditions, where Ni and

Fe transmute to radioisotopes of long half-life. After the workshop, F. W

Mann and D. G. Doran reexamined some of their assumptions in this calculation; we recently learned that their final results are in closer agreement

with the results of Okada.

The effect of Mo additions on 316LN-type alloys was discussed bv P T Purtscher (NBS) . Additions of Mo increase 4-K yield strength more than

additions of Ni, but Mo decreases the 4-K fracture toughness. Thus *:.<• quality index (QI), the product of yield strength and fracture toughne -. -. not changed appreciably. In contrast, additions of Ni improve the frar--ir<-

367
toughness and also the QI because yield strength is not strongly affected.

N. J. Simon (NBS) also discussed 316LN alloy modification by presenting

predictive equations that quantify the effects of Ni, yield strength, and

inclusion spacing upon the 4-K fracture toughness. The work described is an extension of the well-established inverse dependence of fracture toughness upon yield strength. Studies of low- temperature intergranular fracture

in binary Fe-Mn alloys were presented by M. Strum (LBL) . Although this type

of fracture is usually attributed to formation of brittle grain-boundary

phases of grain-boundary precipitation, the work showed that low grainboundary cohesion and heterogeneous slip distribution were more plausible

explanations

.

Improvement of Incaloy 905 by chemistry modification was described

by R„ Ballinger (MIT) . This alloy has a low coefficient of thermal expansion that enables it to be used as a conductor sheath with Nb 3 Sn, a strainsensitive superconductor. Improvements in microstructural stability enable

the modified 905 alloy to withstand the high- temperature , long- duration

thermal processing of the superconductor. However, as A. Nyilas (KFK) indicated, in some designs of cable in conduit magnets long sections may be

laser-welded and the conductor may even be added after it has been thermally

processed.

A study of micromechanical aspects of fracture at 4 K using acoustic

emission techniques was presented by H. Takahashi (Tohoku University)

.

Significant levels of acoustic emission were observed during fracture tests

of three different alloys systems. Internal temperature rises of up to

53 K were reported during discontinuous yielding. The session ended with

a report on a new Al— Li alloy, 2090-T8, by A. Sunwoo (LBL). This alloy is

important because, unlike more common Al alloys, its fracture toughness

increases with decreasing temperature. Progress in welding this alloy was described.

SESSION IV. LARGE-SCALE TOKAMAK AND MIRROR FUSION ENERGY DEVICES:

MATERIALS AND WELDING

Y. Kohsaka (Nippon Kokan K.K.) presented results that showed a change

in the number of serrations (per unit strain) in the s tress - strain curves

of 22Mn—13Cr—7 . 5Ni steel with an increase in the strain rate of the test.

(Discontinuous plastic, or serrated, flow is observed in tensile testing at

4 K when localized adiabatic heating results from a very low specific heat.)

Measured tensile and yield strengths also were affected by strain rate.

Strain rates below 4 x 10' 3 s* 1 did not decrease the measured strengths.

Previous talks by P. T. Purtscher and N. J. Simon had presented the

quantitative effects of Ni and inclusion spacing upon fracture toughness

of 316 LN austenitic base alloys at 4 K. T. A. Siewert (NBS) extended this

work to 4-K weld data, showing that equations with similar coefficients

apply to the weld data and that a variation in the equation improved the

overall fit. Fracture surface dimple density was also correlated with

Charpy V-notch absorbed energy and inclusion spacing.

368
H. Takahashi (Tohoku University) reported results of fracture toughness

testing on 310S and its weldment with and without a high-compliance loading

system. The instability point tended to move close to the maximum load when

a high-compliance spring was used to connect the specimen to the tensile

machine. The 4-K serrations in the stress—strain curve were strongly dependent upon the spring constants.

J. W. Chan (LBL) presented results of gas -tungsten- arc and autogenous electron-beam welds on a high -manganese alloy, 18Mn-16Cr-5Ni-0 . 22N.

Strength—toughness combinations at 4 K were lower than that of the base

metal, owing in part to volatilization of interstitial nitrogen in the

fusion zone. Alloying additions of N, Ni, and Mo improved the properties.

C. McCowan (NBS) discussed tests on weld metals of varying Mn, Mo, and N

compositions to determine the effects of these alloying additions on 4-K

yield strength and 77-K Charpy lateral expansion. The 4-K mechanical properties of several of the new austenitic steels developed in Japan under the

auspices of the Japan Atomic Energy Research Institute (JAERI) were discussed by K. Yoshida. Some thick-plate weld results matched the base-plate

results well, with fracture toughness [Kj c (J)] equal to 200 MPa/m. A status

report on the Demonstration Poloidal Coil (DPC) , a cable- in-conduit design, was also presented.

B. Strauss (Teledyne Engineering Services) described some of the

problems with testing methodology at cryogenic temperatures that will be

investigated in a project sponsored by the Material Properties Council.

Problems, such as adiabatic heating, that affect allowable strain rates

will be addressed for both metals and composites.

A discussion of research needs for superconducting magnets concluded

this session. An scheme for optimizing both the interlaminar shear strength

(or some other fracture parameter) and the allowable radiation dose for com- posites, analogous to the quality index (QI) for alloys and weldments, was discussed. Returning to consideration of alloys, V. Der (OFE) presented

some problems involved in converting data to design allowables. For ex- ample, an allowable stress level of 1/2 ou does not provide the usual safe tv

margin if Cy and cr u are quite close, as is the case for some new alloys and

for cold-worked copper. The different approaches taken in different coun- tries to resolve conductor sheath problems were also discussed. The United

States has invested in modifying and testing alloys to determine compatibility with high- temperature superconductor processing and thermal contraction, whereas the European NET project hopes to solve this problem bv d<- s

strategies. Problems of testing thin plates were also considered.

369

UNITED STATES -JAPAN

COOPERATIVE PROGRAM

371

UNITED STATiS- JAPAN COOPERATIVE PROGRAM

PRINCIPAL INVESTIGATOR: R. L. Tobler

CONTRIBUTING ORGANIZATIONS: National Bureau of Standards,

Tohoku University,

Japan Atomic Energy Research Institute,

Kobe Steel,

National Research Institute for Metals

OBJECTIVES

• Preparation of a draft standard for tensile testing of structural alloys

at 4 K

• Preparation of a draft standard for fracture testing at 4 K

• Administration of a round-robin test program to support the standards

development

• Administration of a one-year research program to evaluate critical

aspects of test standardization at 4 K

• Revision of the draft standards to include results of the one -year

research program

PROGRAM HIGHLIGHTS

• Cryogenic materials scientists were interviewed at government, industrial, and academic institutions in Japan. Their experience and advice

was helpful in preparing the drafts of the test procedure standards.

• Three drafts of a proposed standard 4-K tensile test procedure were

written and submitted for review at Japanese domestic and United Star.esJapan workshops. The fourth revision is now approved and has been

submitted to the American Society for Testing and Materials.

• Three drafts of a proposed standard for 4-K fracture toughness testing

were written and reviewed at Japanese domestic workshops. The fourth

revision is ready for international review.

• Round-robin tensile and fracture tests were conducted on a Fe— 1 3Cr-._

alloy. A summary of results is scheduled for presentation at the 198”

International Cryogenic Materials Conference (ICMC) and will be submitto

for publication in the proceedings.

• An experimental research program was completed by participant::; at Nr four Japanese laboratories. From 70 Jj c fracture toughness tests, th<- effects of specimen size, side grooves, test speed, and fatigue cra< on the test procedure were determined. Five papers containing *:.<• results will be presented at the 1987 ICMC and submitted for public. r . in the proceedings

373
• New methods of fracture characterization were explored, including the use

of acoustic emission and key-curve analysis for fracture toughness determination. The advantages and disadvantages of these methodologies are summarized in two of the research papers that will be submitted to ICMC.

• Four workshops were an integral part of the program:

Standardization of Fracture Toughness Testing of Low-Temperature

Structural Materials at Tohoku University, Sendai, Japan in

March 1986

Low- Temperature Structural Materials and Standards at Japan

Atomic Energy Research Institute in Tokai-mura, Japan in

Augus t 1986

DOE—JAERI Low Temperature Structural Materials and Standards

Workshop in Reno, Nevada in October 1986

United States—Japan Cooperative Research Program Review at

Tohoku University in January 1987

374
UNITED STATES-JAPAN

LOW- TEMPERATURE STRUCTURAL MATERIALS AND STANDARDS WORKSHOP

Airport Plaza Hotel

Reno , Nevada

October 9—10, 1986

The results of the United States—Japan collaboration on structural materials development and test-method standardization were presented and discussed at this meeting of program

participants. Research goals to be accomplished before the

next meeting, at the Japan Atomic Energy Research Institute in

1988, were formalized.

375

UNITED STATES-JAPAN

LOW TEMPERATURE STRUCTURAL MATERIALS AND STANDARDS WORKSHOP

Airport Plaza Hotel

Reno , Nevada

October 9—10, 1986

PROGRAM

October 9

SESSION I. OVERVIEW

9:00 — V. K. Der (DOE) and S. Shimamoto (JAERI) , cochairpersons

10:30 a.m.

A. Introduction and Welcome

B. JAERI Program Overview,

S. Shimamoto and K. Yoshida

C. DOE Program Overview,

V. K. Der

D. NBS Program Overview,

R. P. Reed

SESSION II. TESTING

11:00 a.m. — T. A. Siewert (NBS) and K. Ishikawa (NRIM) , cochairpersons

12:10 p.m.

A. Fracture, Adiabatic Heating, and Acoustic Emission of AusteniteStainless Steels at Liquid Helium Temperature,

H. Takahashi and T. Shoj i (Tohoku University)

B. Dynamic Yield Stress and Fracture Toughness of Some Material.-. • Low Temperature,

K. Kishida, T. Kataoka, and M. Nakano (Osaka University)

C. Evaluation of Tensile and Fracture Properties of Austenitic ir

Alloys in Liquid Helium,

K. Ishikawa, T. Ogata, T. Yuri, and K. Nagai (NRIM)

D. Construction of 100 Ton Cryogenic Testing Machine and 7- Results of Newly Developed Stainless Steels with This M.. : 7 : • at 4 K,

S. Shimamoto, H. Nakajima, K. Yoshida, and M. Oshikiri J AH:- :

377
SESSION II. TESTING, continued

1:30 -

2:30 p .m.

E. Strategy for the Development of the Data Base for the RadiationResistant Superconductive Magnet Materials,

T. Okada and S. Nishijima (Osaka University)

F. Research on Development of Low Temperature Test Standards and

Property Measurements

,

R. P. Reed (NBS)

G. Standardizing Nonmetallic Composite Materials for Cryogenic

Applications

,

M. B. Kasen (NBS)

SESSION III. MATERIALS

2:30 — L. T. Summers (LLNL) and T. Okada (Osaka University),

5:30 p.m. Cochairpersons

A. Fatigue Properties of 22Mn Steel in Cryogenic Temperature,

M. Shimada, S. Tone, T. Mizoguchi, and T. Horiuchi (Kobe Steel)

B. Low Temperature Mechanical Properties of 12Cr—12Ni—lOMn—5Mo Steel

and its Weld,

R. Miura and J. Ishizaka (Japan Steel Works)

C. Change in the Cryogenic Properties of the Welded Joint of

V-bearing Austenitic Stainless with Nb 3 Sn Precipitation Heat

Treatment

,

K. Nohara (Kawasaki Steel)

D. Mechanical Properties of Welded Joints in Austenitic Stainless

Steels at 4 K,

C. Ouchi and K. Kohsaka (Nippon Kokan K.K.)

E. Mechanical Properties of Fully Austenitic Weld Deposits for

Cryogenic Structure,

T. Matsumoto (Hitachi)

F. Research on Structural Steels for High-Field Superconducting

Magnets

,

J. W. Morris, Jr. (LBL)

G. SMA and FCAW Compositions for Magnetic Fusion Superconducting

Magnet Cases,

D. J. Kotecki (Teledyne McKay)

H. A Fully Austenitic Stainless Steel Weldment for Cryogenic Service,

F. S. Babish (Sandvik) and T. A. Siewert (NBS)

378
October 10

SESSION IV. DESIGN EXPERIENCE AND STANDARDS

9:00 — H. Becker (MIT) and K. Yoshida (JAERI) , cochairpersons

10:00 a.m.

A. Mechanical Design Guideline of the Superconducting Coils for the

Fusion,

K. Yoshida, K. Koizumi, H. Nakajima, and S. Shimamoto (JAERI)

B. Thoughts on Structural Design Standards for Magnet Casings,

H. Becker (MIT)

C. MFTF-B: Experience with Structural Materials and Operating

Stresses

E. N. C. Dalder, J. 0. Myall, R. M. Scanlon, and

J. P. Zbasnik (LLNL)

SESSION V. OPEN DISCUSSION:

in AA JOINT DEVELOPMENT OF TESTING STANDARDS

10:20 a.m. —

12 ; 40 p . m.

A. Progress Towards 4-K Tensile and Fracture Toughness Testing

S tandardizat ion

H. Takahashi, T. Sho i i , and R. L. Tobler (Tohoku University

and NBS)

B. Round-Robin Test Results for Tensile and Fracture Toughness

Testing Standardization,

H. Takahashi, K. Ishikawa, M. Shimada, T. Matsumoto,

S. Shimamoto, and H. Nakajima (Tohoku Univesity, NRIM,

Kobe Steel, Hitachi, and JAERI)

C. Evaluation of Low Temperature Mechanical Properties of

Fe—22Mn—13Cr—5Ni Austenitic Steel,

L. Summers (LLNL)

D. Round Robin Test Results on 316LN Stainless Steels in the Uni-.ed

States

P. T. Purtscher (NBS)

12:40 p.m. SESSION VI. SUMMARY

379

UNITED STATES-JAPAN

LOW- TEMPERATURE STRUCTURAL MATERIALS AND STANDARDS WORKSHOP

SUMMARY OF TECHNICAL PRESENTATIONS

SESSION I. OVERVIEW SESSION

S. Shimamoto presented an overview of the Japan Atomic Energy Research

Institute (JAERI) program on developing new cryogenic structural steels

in collaboration with Japanese steel companies. Electron-beam and gastungsten-arc welding of thick plate is progressing, with expectations that

90% of the base metal properties will be retained and that the Kj c fracture

toughness values of the weldments at 4 K will be 200 MPaVm. Nine organizations in Japan are capable of cryogenic Kj c measurements, and round-robin

testing has begun. R. L. Tobler (NBS) is helping to standardize cryogenic

test procedures. Construction of the Fusion Experimental Reactor (FER) is

expected to start in 1992. Meanwhile, experience is being acquired in the

Demonstration Poloidal Coils (DPC) program, where both Nb-Ti and Nb 3 Sn will

be used as conductor materials.

V. Der (OFE) described current structural alloy research at NBS, LLNL,

LBL, and MIT, summarizing many of the details reported on in the preceding

Cryogenic Structural Materials Workshop). Test methods and standards development were described for both alloys and nonmetallics . Progress in

round-robin measurements was noted, as were exchanges of personnel and information. Additional work and collaboration is needed in the development

of design guidelines and standards.

The low- temperature structural materials research at NBS was reviewed

by R. P. Reed. Study of 316LN-based alloys from a number of sources led t predictive equations for their 4-K yield strength and fracture toughness,

indicating that the U.S. goals for these properties can be met through cor. - trol of chemistry, grain size, and inclusion spacing. Weldment propert:--; are still challenging. Progress in the development of standard nonme ta

specimens and test methods was described.

SESSION II. TESTING SESSION

H. Takahashi (Tohoku University) discussed a key-curve method o: : . , ture toughness measurements, a simpler and cheaper method of 4-K frac;

toughness evaluation. Success of this method depends upon the existence

a correlation between Jj c and the slope dJ/da for the alloys of inter*-?,:

Further evaluation is required. Strain rate effects on fracture 'on,-:.: . , and yield strength measurements were presented by K. Kishida (Osaka

versity) . In the static limit, the curve of Kj c versus temperature

above the dynamic testing results. K. Yoshida (JAERI) present'--! r. of studies on specimen size effects on tensile properties. .‘Jo

• found for diameter variations from 24 to 7 mm. Work on Kj c specimen

effects will be carried out this fall in collaboration with R. h

381
A newly constructed 1000 -kN cryogenic testing machine will be used in this

work.

The problem of radiation degradation and radioactivation of magnet

alloys was addressed by T. Okada (Osaka University). The results given

for the time for radiation to subside to acceptable levels for maintenance

operations differed from those presented by F. Mann earlier in the Cryogenic

Structural Materials Workshop. Thii' difference is due to the different

radiation energy spectra assumed in the two calculations. Mann used the

Starfire configuration, which assumed a soft spectrum at the magnet. Okada

used a spectrum that assumed 10% hard (14 MEV) neutrons so that problems

with Ni and Fe activation were similar to those encountered at the first

wall; his calculation used a flux of ~10 10 n/cm2 s. The Starfire-based

calculation has since been revised (see discussion in section III of the

Ninth Annual Cryogenic Structural Materials Workshop Summary.)

R. P. Reed (NBS) discussed strain-rate effects on measured tensile

properties of alloys at 4 K. Effects on flow strength were minimal at 4 K.

In discontinuous yielding, the measured specimen temperature rise (AT) of

about 50 K for 304L at 4 K was consistent with results of a model in which

no heat was conducted away from the deformed region (adiabatic heating)

.

Calculations for 316LN tests at 4 K showed that the observed temperature

rise with strain rate increase coincided with the transition from nucleate

to film boiling heat transfer in the liquid helium, which results in an increased AT between the solid and liquid.

M. B. Kasen (NBS) discussed a coding scheme for identifying composite

materials in a standardized manner. This would allow for nonmetallic

specification in the way that the AISI specifications for stainless steel

identify those materials. Such a system has not been available for

nonme tallies

.

SESSION III. MATERIALS SESSION

Fatigue tests on 22Mn steel at 4 K were reported by M. Shimada (Kobe

Steel). Surface temperature rise was studied as a function of stress and

frequency. Serrated yield was reported as a function of strain rate and

compared with that of 304LN. Strain- and stress -controlled fatigue life,

fatigue crack growth data, and threshold fatigue data were reported.

Mechanical properties of 12Cr- 12Ni- lOMn- 5Mo steel and weldments were

reported by R. Miura (Japan Steel Works) . Electroslag remelted base

material and GTA weldments exceeded the minimum JAERI criteria for yield

strength and toughness. A new 316- type austenitic stainless steel with

low-C, high-N, and V additions was found to retain favorable 4-K properties

after the type of heat treatment required by the Nb 3 Sn superconductor. This

work at the Kawasaki Steel Corporation was described by K. Nohara. Plate

was prepared by vacuum induction melting and EB welding was utilized. The

gas- tungsten- arc and electron-beam welding of Mn-Cr and Ni-Cr austenitic

stainless steels that met the JAERI criteria was reported by Y. Kohsaka

(Nippon Kokan K.K). Satisfactory 4-K test results were obtained. Fully

382
austenitic weld deposits were studied by T. Matsumoto of Hitachi. Effects

of Mo, C, and C+N on fracture toughness were reported. No delta ferrite was found in the weld microstructure, magnetic permeability was not significantly larger than for base material, and microfissuring was not observed.

J. W. Morris (LBL) reported results of fundamental studies on fracture

mechanisms that could explain the favorable trend of fracture toughness

of aluminum alloy 2090 as temperature is decreased. As the temperature

decreases, the fracture mode does not change, but elongation increases

because da/de increases

D. J. Kotecki (Teledyne McKay) described welding filler metal development

work for the fabrication of MFTF-B at LLNL. Lateral expansion was improved

for compositions to the left of the 0 ferrite number line on the DeLong diagram, and this criteria proved suitable for the selection of filler material

for flat welding positions. However, a ferrite number of 3 was necessary

for vertical-up position welding. Satisfactory Kj c properties were obtained

with this composition change.

Yield strength and Kj c(J) at 4 K for 18Cr-20Ni-5Mn-0 . 16N austenitic

weld metal were reported by T. A. Siewert (NBS) . The properties obtained

from a GMA weld process were comparable to base metal properties, and SEM

analysis showed ductile failure by microvoid coalescence.

SESSION IV. DESIGN EXPERIENCE AND STANDARDS SESSION

A document setting forth mechanical design guidelines for superconducting magnets for a tokamak fusion machine was discussed by K. Yoshida

(JAERI) . Tests are required for each production batch of alloys and non- metallics because the existing data base is not sufficient. Design allowables, fabrication, and inspection are also covered. The guideline is base i on the 1986 ASME Boiler and Pressure Vessel Code. Yoshida concluded Chat

it was impractical at present to use high-performance cryogenic steels in i tokamak machine, and that use of these new steels to improve current dens:*:

would be the most effective application.

H. Becker (MIT) presented a draft "Structural Design Basis." This dc . ment suggests that since it may be impossible to eliminate all cracks or r possibility of crack initiation, it may be necessary to assume crack exi , tence as a basis for design and to provide for remotely controlled ins:

tion during service.

J. Zbasnik (LLNL) described MFTF-B experience with structural mar.-r i and operational stresses. Cryogenic testing of the 304LN production pla~.- was required. Tensile properties usually exceeded the specifications

Properties of the qualified welding electrodes were similar to aust.-ni~i

plate properties. Allowable stresses depended on types of loading and : im- probability of occurrence. The magnets were instrumented with strait:

to test agreement with predicted values. In all cases, measured x-

were lower than or in good agreement with design values, which valid. if.-- design approach.

383
SESSION V. JOINT DEVELOPMENT OF TESTING STANDARDS

The status of cooperative 4-K tensile and fracture toughness testing

standardization was described by H. Takahashi. A draft of a proposed

standard was distributed earlier this year, and initial round-robin tests

have been completed. Supportive research on a number of aspects of testing

is also in progress. R. L. Tobler (NBS) has been participating in this

program in Japan, S. Shimamoto (JAERI) set forth the results of 4-K roundrobin tests on a 22Mn—2Cr—5Ni steel carried out at five laboratories. Both

tensile and fracture toughness tests showed good correspondence. Some

research questions for fracture toughness test standardization have been

answered; others are still under investigation. Further round-robin tests on 316LN stainless steel and a new Japanese cryogenic steel are planned.

L. Summers gave the results of Kj c (J) tests at LLNL on the 22Mn steel;

more scatter was found in results from specimens of TL orientation than

of LT orientation, but the results were in general agreement with the

Japanese results (TL orientation) . A comparison of Japanese and NBS test

results on the 22Mn steel specimens was presented by P. T. Purtscher (NBS).

Especially good agreement was noted for yield strength results. Fracture

toughness testing procedures were detailed. The TL Kj c (J) results from NBS,

with a mean of 156 Ma^m, are lower than most of the Japanese test results;

LT results were in better agreement. A change in Kj c (J) test procedure

is that fatigue cracking is now carried out at 77 K rather than at room

temperature

.

SESSION VI. SUMMARY SESSION

The following action items were agreed upon and will be carried out

before the next United States—Japan Workshop, which will be held at JAERI

in the fall of 1988.

Test Methods and Standards

o Participants from both the United States and Japan will continue

research on test methods and standards. R. L. Tobler will continue

collaboration on Kj c (J) standards.

o Japan will send machined specimens to the United States for a second

round-robin test series. The U.S. test results will be reported prior

to the next workshop.

o The U.S. participants will draft a summary report on the first roundrobin test results.

o The round-robin testing on 316LN specimens will be completed.

o A joint paper on test standards will be drafted: tensile tests will

be addressed in one part, Kj c (J) in another part of this paper.

o A plan will be developed for comparing test methods on strengths of

composites

.

384
Materials

A. Alloys

o At JAERI , tests of large specimens will determine the true Kj c instead of

Jj c and may enable future large specimen tests.

o A test matrix leading to a data base will be composed (this should determine whether observed variability arises from the microstructure, metastability, or mechanistic causes.)

o A discussion of the factors and parameters that should be considered in

the formulation of a test matrix for an adequate data base will initiate

with a letter from J. Morris (LBL) to S. Shimamoto (JAERI). This letter

will urge the expansion of the program to include fundamental metallurgical studies of relevant alloys by participation of one or more Japanese

university professors of metallurgy.

B. Weldments

o Japanese participants will provide a representative plate alloy and a welded thick plate to the U.S. participants for conducting welding tests.

The test specimens used and any results obtained will be shared with

Japanese participants. The U.S. participants will explore possible welding processes and electrodes.

C. Nonme tallies

o Materials samples will be exchanged.

o Jointly, Japanese and U.S. participants will develop a test matrix for

characterization of model commercial materials and will proceed with the

program after agreement on the test matrix.

D. Design Standards and Guidelines

o Data on past magnet design and experience leading to design changes will

be acquired.

o Reports on magnet failure (structural failures) will be collected.

o Data on failure and failure prevention (fracture control) will be

collected.

o U.S. and Japanese participants will draft a preliminary recommended b *; i for structural design. Specialists from each country will meet m dis- cuss the specifics of the design basis.

385



JAPAN—U . S . COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

FOURTH DRAFT

MARCH 1987

NOTE: This is a working document of a standard being developed by the

above committee for eventual submission to ASTM; it is intended

for committee review purposes, but not for general distribution

at this time.

Proposed Standard Method for

TENSILE TESTING OF STRUCTURAL ALLOYS AT LIQUID HELIUM TEMPERATURE

1 . Scope

1.1 This standard describes the procedures to be used for the tensile

testing of metals in liquid helium [4 or 4.2 K (—452°F)]. The format is

similar to that of other ASTM tensile test standards, but the contents include modifications for cryogenic testing, which requires special apparatur

smaller specimen size, and concern for serrated yielding, adiabatic heatir..:

and strain rate effects.

1.2 To conduct a tensile test by this standard, the specimen in a cry

stat is fully submerged in normal liquid helium (He I) and tested using

crosshead displacement control at a moderate strain rate. Tests that us^

load control or high strain rates are not considered in this standard

1.3 This standard details methods for the measurement of vield stre: J tensile strength, elongation, and reduction of area. The determination

the elastic modulus is treated in Method E 111.

1.4 Values stated in SI units are treated as primary; values sta:> : U.S. customary units are treated as secondary.

1.5 This standard does not address the issue of safety associated

cryogenic testing. It is the user's responsibility to learn and obser-^

appropriate safety practices. Safety guidelines for handling liquid

and other cryogens are available elsewhere (1)

NOTE 1 —The boiling point of normal liquid helium (He I) a' sea level is 4.2 K. This value decreases slightly with geographic elevation; it is 4.0 K at NBS in Boulder, Colorado.

1677 m (5500 ft) above sea level. In this standard, the test

temperature is nominally designated 4 K.

387
2. Applicable Documents

2 .

1

ASTM Standards:

A 370 Methods and Definitions for Mechanical Testing of Steel Products

(ANNUAL BOOK OF ASTM STANDARDS, Vol. 01.04)

E 4 Practices for Load Verification of Testing Machines

(ANNUAL BOOK OF ASTM STANDARDS, Vols. 03.01, 04.02, 07.01, and 08.03)

E 6 Definition of Terms Relating to Methods of Mechanical Testing

(ANNUAL BOOK OF ASTM STANDARDS, Vols. 03.01 and 08.03)

E 8 Methods for Tension Testing of Metallic Materials

(ANNUAL BOOK OF ASTM STANDARDS, Vols. 01.02, 02.01, 02.02, 02.03,

and 03.01)

E 8M Methods for Tension Testing of Metallic Materials, Metric

(ANNUAL BOOK OF ASTM STANDARDS, Vol. 03.01)

E 29 Recommended Practice for Indicating Which Places of Figures Are

to Be Considered Significant in Specified Limiting Values

(ANNUAL BOOK OF ASTM STANDARDS, Vols. 02.03, 03.01, 03.03, 03.05,

and 14.02)

E 83 Practice for Verification and Classification of Extensometers

(ANNUAL BOOK OF ASTM STANDARDS, Vol. 03.01)

E 111 Method for Young's Modulus, Tangent Modulus, and Chord Modulus

(ANNUAL BOOK OF ASTM STANDARDS, Vol. 03.01)

3. Definitions

3.1 The definitions of terms relating to tension testing that appear in

ASTM Standard E 6 shall apply here. The following definitions also apply:

3.1.1 tensile crvostat—a test apparatus for applying loads to test

specimens in cryogenic environments. A schematic illustration is shown in

Fig. 1.

3.1.2 Dewar—a vacuum- insulated container for cryogenic fluids.

3.1.3 adiabatic heating—the internal heating of a specimen resulting

from tensile testing under conditions such that the heat generated by plastic work cannot be quickly dissipated to the surrounding cryogen.

3.1.4 reduced section—section in the central portion of the specimen,

which has a cross section smaller than the gripped ends.

3.1.5 length of the reduced section—the distance between the tangent

points of the fillets that bound the reduced section.

388
3.1.6 adjusted length of the reduced section—the length of the reduced

section plus an amount calculated to compensate for strain in the fillet

region.

3.1.7 gage length—the original distance between gage marks made on the

specimen for determining elongation after fracture.

3.1.8 axial strain—the average of the strain measured on opposite

sides of the specimen at equal distances from the specimen axis.

3.1.9 bending strain—the difference between the strain at the surface

of the specimen and the axial strain. The bending strain varies from point

to point around and along the reduced section of the specimen.

3.1.10 maximum bending strain—the largest value of bending strain in

the reduced section of the specimen. It is calculated from strains measure:,

at three circumferential positions, at each of two different longitudinal

positions

4. Significance and Use

4.1 In general, tensile tests provide information on the strength and

ductility of materials under uniaxial tensile stresses. This information

may be useful for alloy development, comparison and selection of materials,

and quality control. Under certain circumstances, the information may also

be useful for design.

4.2 At 4 K, the load-time and load-deflection records for metals testin displacement control are serrated (2). Serrations are formed by repe:'

bursts of unstable plastic flow, followed by arrests. The unstable plasti

flow (or discontinuous yielding) is a free-running process that occurs

locally with sizable heat evolution at higher than nominal rates of strair

Typical stress—strain curves for an austenitic stainless steel with serra:

yielding at 4 K are shown in Fig. 2.

4.3 Constant specimen temperature cannot be maintained at all times

during tests in a liquid helium environment. Owing to adiabatic heat:.

specimen temperature temporarily rises above that of the cryogen at -* during the discontinuous yielding events (see Fig. 2). This behavior var:

with specimen size and testing speed, but altering the mechanical

ables cannot eliminate the periods of internal heating (3). Therefor--

sile property measurements of alloys in liquid helium (especially

strength, elongation, and reduction of area) lack the full significant.-

property measurements at higher temperatures where discontinuous y i * : does not occur.

4.4 At 4 K, the stress—strain behavior of a material during uns* -I: •• plastic deformation depends on whether load control or displacement

is used (4). Crosshead displacement control is specified because -h«-

389
of this standard is material characterization by conventional methods. This

limitation must be taken into account when data are used for structural design in applications when actual circumstances may approach load-controlled

conditions

.

5 . Apparatus

5.1 Test Machines—Machines used for tensile testing shall meet the

requirements of Practices E 4 regarding verification of load accuracy.

5.2 System Design—Alloy strengths often double or triple between room

temperature and 4 K. For the same specimen geometry, higher loads must be

applied to the cryostat, test specimen, load train members, and grips at

cryogenic temperatures. For most conventional test machines, which have a maximum load capacity of 10 tons, it is recommended that the apparatus be

designed to accommodate one of the small specimens cited in section 7.2.2

of this standard.

5.3 Construction Materials—Many construction materials, including the

vast majority of ferritic steels, are brittle at 4 K. To resist embrittlement, the grips and other load- train members must be fabricated using

strong, tough, cryogenic alloys. Austenitic stainless steels (AISI 304LN)

,

maraging steels (200, 250, or 300 grades, with nickel plating to prevent

rust), and extra- low- inters titial (ELI) grade titanium alloys (Ti—6Al—4V

and Ti—5A1—2 . 5Sn) have been used, with proper design, to fabricate grips,

pull rods, and cryostat frames.

5 . 4 Alignment :

5.4.1

Single - Specimen Apparatus—To avoid bending strains in the specimen, proper system alignment is essential. For a conventional singlespecimen test apparatus, the machine and grips shall be capable of loading

a precisely machined specimen so that the maximum bending strain does not

exceed 10% of the axial strain. This calculation is based on the strain

readings taken at zero load and at the highest load for which the machine

is being qualified.

NOTE 2—This requirement will minimize contributions from

the test apparatus to the bending strain. Tests performed

with a qualified apparatus may still vary in percent bending

strain owing to small variations in the specimens.

5.4.2

Multiple-Specimen Apparatus—For this type of cryostat, the

alignment depends on the type of fixtures used. The maximum bending strain

shall be measured and reported.

5.4.3

Axialitv Tests—The testing apparatus may be qualified by

axiality measurements at room temperature and at 4 K. To perform axiality

tests of the apparatus, the specimen form should be the same as that used

during cryogenic tests,

390
the specimen concentricity should be as nearly perfect as possible, and no plastic strains should occur in the reduced section. In some cases, this

may necessitate the use of a relatively stiff, high-strength calibration

specimen.

5.4.3. 1 For round bar specimens, the maximum bending strain, as defined

in 3.1.10, is calculated from the strains measured at three circumferential

positions, at each of two different longitudinal positions. The strains are measured with three electrical-resistance strain gages equally spaced arounc

the reduced section of the specimen. The two longitudinal positions should

be as far apart as possible, but not closer than one diameter to a fillet.

5. 4. 3. 2 For specimens of rectangular cross section, strain may be

measured at the center of each of the four sides, or in the case of thin

strips, near the outer edges of each of the two broad sides.

5. 4. 3.To eliminate the effect of specimen bias, repeat the axiality

measurements with the specimen rotated 180 degrees, but with the grips and

pull rods retained in their original positions. The maximum bending strain

and strain at the specimen axis are then calculated as the average of the

two readings at the same position relative to the machine.

5.4.4 Strain Averaging Technique—Nonaxiality of loading is usually

sufficient to introduce errors in tensile tests at small strains when strain

is measured on only one side of the specimen. To rectify this, two extensometers or strain gages may be attached to opposite sides of the specimen.

The reported strain should be the average of the strains on both sides.

5.5 Gripping Mechanisms—The choice of gripping mechanism to be used

at 4 K is influenced by specimen type. Any suitable mechanism described ir. Methods E 8 and E 8M may be used.

5.6 Dimension-Measuring Devices—Micrometers and devices used for mea- suring the dimensions of specimens shall be accurate and precise to at lea.

one-half the smallest unit to which a given dimension must be measured.

5 . 7 Cryostats and Support Apparatus :

5.7.1 Cryostats—A cryostat capable of retaining liquid helium is pre

requisite. In general, cryostat load frames for existing test machines t.

be custom-built, but they may accommodate commercially available Dewars

cryostat may employ adjustable load columns to facilitate alignment.

practical designs, including turret-disc designs for multip le - spec imen testing with a single cooling, are discussed in the literature (5-9)

5.7.2 Dewars—Stainless steel Dewars are recommended because the-.- ,i;. safer than glass Dewars and less expensive than fiberglass Dewars

helium Dewar (see Fig. 1) is usually sufficient for short-term tests.

double-Dewar arrangement in which an outer Dewar of liquid nitrogen ; the inner Dewar of liquid helium is also possible.

391
5.7.3 Ancillary Equipment—Dewars and transfer lines for liquid helium

must be vacuum insulated. Vacuum pumps, pressurized gas, and liquid nitrogen facilities are therefore required. After testing, the helium may be

released to the atmosphere, recycled as a gas, or reliquefied. Recycling

or reliquefaction requires large investments in purification and support

systems

.

5 . 8 Temperature Maintenance and Liquid Level Indicators :

5.8.1 Thermocouples—The intended temperature of 4 K is ensured by

maintaining a liquid helium environment. A thermocouple to measure the

specimen temperature is not required for routine tests.

5.8.2 Indicators—Although a thermocouple attached to the specimen is

not necessary, an indicator or meter is required to ensure that the specimen

remains fully submerged for the duration of testing. On—off indicators of

the carbon-resistor type may be used to verify that the liquid level always

exceeds some reference point above the specimen, or the liquid level may

be continuously monitored with superconducting wire sensors of appropriate

lengths positioned vertically inside the cryostat.

5 . 9 Strain Gages :

5.9.1 Selection—Strain-gage films bonded directly to the specimen

surface may be used to measure strain at 4 K (10). A satisfactory combination of gage active element, backing material, and bonding agent should

be selected on the basis of experience and manufacturer's recommendations.

One common choice is a Ni—Cr-alloy gage with a temperature -compensated

active element (7,11).

5.9.2 Characteristics—Strain gages are typically wired to a dummy

bridge using a three-wire temperature-compensating hookup, like that shown

in Fig. 3. The gage resistance is typically 120 or 350 0. A low excitation

voltage of about 1 to 2 V is recommended at 4 K to reduce Joule heating.

Typical full-scale operating ranges are 1% at room temperature and 2% at

4 K.

5.9.3 Cal ibration—Strain gage calibration at room temperature requires

calibration of the electronics to changes in resistance. Calibration at 4 K

requires consideration of the temperature dependence of the resistance in

the electronic calibration. Gage factors as a function of temperature are provided by the manufacturers and in published research.

5.10

Clip-Gage Extensome ters :

5.10.1 Types—Detachable clip-gage extensometers for use at 4 K may be

built or purchased. An example is the beam gage, which uses four straingage films bonded in a Wheatstone bridge arrangement (11). Extension within

the specimen gage length is sensed by the extensometer, which is clipped to

retaining pins that are fixed to the specimen reduced section.

392
5.10.2 Characteristics—To measure the 0.2% offset yield strength,

a Class B—2 extensometer , as identified in Method E 83, may be used. The

extensometer shall meet the sensitivity and accuracy requirements of Method

E 83 and shall be tested to ensure accuracy at 4 K. Whenever possible, the

extensometer should be mounted directly to the specimen reduced section.

5.10.3 Calibration—Extensometers should be calibrated at room temperature and at 4 K with a suitable device. For calibrations at 4 K, a micrometer with vertical extension tubes can be used with the extensometer (s) mounted at the lower end and immersed in liquid helium. Once the

calibration is known and proven to be accurate, linear, and reproducible,

room- temperature checks may be performed prior to each test series for indirect verification of the calibration at 4 K. Direct calibration at 4 K,

however, must be performed periodically, when damage is suspected or repairs

have been made.

5.11 Capacitance Extensometers—Extensometers that use capacitance

measurement to monitor strain may be used (10). The type with overlapping

concentric cylinders has an extended strain range, an output that is linear

with displacement, and an adjustable sensitivity. The type with parallel

plates has high sensitivity, but the output must be compensated for the

hyperbolic dependence of the capacitance on displacement.

6 . Sampling

6.1 Stock Condition—Samples for tensile testing should be taken from

the material in its final condition to ensure that the properties measured

are representative of the end product. Allowance should be made for any

superficial effects introduced by cutting operations.

6.2 Location—Specimens should be cut from locations thought to be mos*

representative of the stock material, realizing that test results for

specimens taken from selected locations of a part or material may not be

representative of the whole. The conventional locations should normally be

used:

6.2.1 For products 40 mm or less in thickness, diameter, or distance

between flats, the location should be at the center.

6.2.2 For products over 40 mm in thickness, diameter, or distance between flats, the location should be midway from the surface to the center

6.3 Specimen Configuration—The choice of specimen size and shape

should be based on the requirements necessary to obtain representative

samples of the material, and on test machine load capability (section

6.4 Orientation—The orientation of the specimen axis relative * • •• principal working directions of the stock shall be specified us in.- notation in section 4 of Methods A 370.

393
7. Test Specimens

7 .

1

General

:

7.1.1 Types and Specifications—Any specimen configuration cited in

Methods E 8 or E 8M may be used. The specifications for dimensions, tolerances, and surface finish are stated in sections 6.1 through 6.17 of those

standards

.

7.1.2 Size—Specimens from sheet or wire products having relatively

small cross-sectional areas can be tested within the load limits of conventional apparatus. Specimens from thick plate or bar products, however, may

need to be machined to a reduced cross-sectional area so the load capacity

of the machine is not exceeded.

7 . 2 Round Bar Specimens :

7.2.1 Standard Room-Temperature Specimen—A 12.5-mm (0.5- in) diameter

round bar specimen is the standard configuration for room- temperature tests

according to Methods E 8 and E 8M. For strong alloys, however, this

specimen requires excessive loads at 4 K. For example, 210 kN is required

to test typical AISI 304LN steel at 4 K, whereas 100 kN or 10 tons is the

limit for most machines.

7.2.2 Standard 4-K Specimens—To meet the load limitations of con- ventional test machines, the round bar specimens in sections 7.2.2. 1 and

7. 2. 2.

2

are recommended as standard for 4-K tests. The required dimensions

and tolerances for these specimens are given in Table 1. Threaded or shouldered ends are common for gripping these specimens, and the requirement of

section 5.4.1 can be met by precise machining.

7.2.2.

1

Standard, small metric specimens . These specimens have a 7-mm

diameter and a gage - length- to-diameter ratio of 5:1.

7. 2. 2.

2

Standard, small U.S. customary specimens . These specimens have

a 6.25-mm (0.25-in) diameter and a gage- length- to-diameter ratio of 4:1.

7.2.3 Alternatives—If the 4-K standard specimens recommended above are not appropriate, other sizes may be selected following the guidelines of

Methods E 8 and E 8M . The proportions of such specimens should be similar

to those of the standard specimens (see Fig. 4 of this standard and Fig. 8 of Methods E 8 and E 8M)

.

7.2.4 Subsize specimens—Special care in fabrication and testing is

required for specimens with diameters less than 6 mm. As the specimen size

is reduced, factors such as machining, surface finish, alignment, and the

number of grains per cross section are of increasing importance.

394
8 . Procedures

8 . 1 Marking and Measuring the Test Specimen :

8.1.1 Gage Length—Gage length marks should be lightly punched,

scribed, or inked at appropriate locations on the reduced section of the

specimen, which is the conventional method. The gage length should normally

be five times the diameter for metric specimens or four times the diameter

for U.S. customary specimens. If another gage length is used, it must be

described in the report. Measure the gage length to the nearest 0.05 mm.

NOTE 3—For metals of low ductility, gage marks punched

or scribed on the reduced section may induce failure at those

locations due to stress concentrations. To avoid this, it may

be possible to coat the reduced section with layout ink, and

then mark the gage length by rotating the specimen in a jig

with knife edges scraping off the ink at the appropriate

intervals. Alternatively, gage marks may be placed on the

specimen shoulders, or the overall length of the specimen may

be measured.

8.1.2 Reduced Section—Measure the length of the reduced section and

the adjusted length of the reduced section, if applicable, to the nearest

0.05 mm.

8.1.3 Cross Section—Measure the cross-sectional area of the reduced

section as specified in section 7.1.1 of Methods E 8 and E 8M.

8.2 Specimen Installation—Install the specimen in the cryostat,

leaving sufficient slack for instrumentation wires so they will not be

stretched or crimped during positioning of the Dewar and subsequent testi:..-

8.3 Seating and Alignment—If the gripping fixture involves loosefitting components, such as spherically seated bearings, prevent friction

or mismatch of the bearing surfaces at 4 K by first checking the seating

and alignment at room temperature. During this process of alignment, the

applied loads should be kept below one-third of the proportional limit of

the test material.

8.4 Cooling Procedure—Remove any moisture from the apparatus pr: :>r to cooling by drying it thoroughly with a pressurized air jet or heat ,• Ice can block cryogenic transfer lines or cause erratic loading behavi r it forms between various parts of the specimen, clip gage, and load tr.v.: Next, position the Dewar and precool the apparatus to 77 K by transfer:::

liquid nitrogen into the cryostat. After attaining equilibrium at ' K

remove all the liquid nitrogen from the cryostat, and transfer liquid -.<•! . into the cryostat until the specimen and grips are fully submerged. 7<-s ti:

may begin after the system has reached thermal equilibrium at 4 K. 7. specimen must remain fully at all times during the test.

395
NOTE 4—The heat transfer characteristics of gaseous helium

are inferior to those of liquid helium; therefore, a liquid

helium environment is required to minimize specimen heating

effects during discontinuous yielding.

8 . 5 Testing Speed :

8.5.1 Rate Control—Owing to adiabatic heating, tensile property

measurements at 4 K can be significantly affected by the testing speed.

Therefore, the test procedure must include a means of measuring and con- trolling the rate of crosshead motion. A nominal strain rate must be

specified, since the actual rate cannot be precisely controlled or maintained when discontinuous yielding occurs. The nominal strain rate is

calculated by dividing the crosshead rate by the length of the reduced

section. Alternatively, a pacing or indicating device may be used to

monitor the strain rate or an average strain rate may be determined by

observing the time required to effect a known increment of strain.

8.5.2 Rate Limit—The nominal strain rate at any time during the test

shall not exceed 1 x 10" 3 s’ 1

. Higher rates may cause excessive specimen

temperature rises and therefore are not satisfactory for determining the

acceptability of materials.

8.5.3 Rate Change—The strain at which discontinuous yielding begins

usually increases with decreasing strain rate. If the first serration for

a given material occurs near 0.2% plastic strain, it may be desirable to

reduce the speed to prevent interference in the measurement of the yield

strength (see Fig. 5). Then, a relatively low strain rate may be used to

determine the yield strength, followed by a higher strain rate to complete

the test. Any convenient crosshead speed may be used up to a stress of one

half the yield strength; after that, the crosshead speed must be such that

the nominal strain rate does not exceed the 1 x 10’ 3 s* 1 limit.

8 . 6 Measurement of Mechanical Properties :

8.6.1 Load—Extens ion Curve Method—To measure the yield strength, a record of load versus extension must be obtained, up to at least 0.2%

plastic strain. The use of a strain measurement device for autographic

recording is recommended. Measure the yield strength via the 0.2% offset

method, following Methods E 8 or E 8M, section 7.3.1. If the 0.2% offset

line intersects the curve at a load drop owing to discontinuous yielding,

then report the highest stress prior to that load drop as the yield

strength (see Fig. 5).

8.6.2 Load—Time Curve Method—Yield strength measurements based on a 0.2% offset procedure applied to load-versus - time curves at 4 K are not

recommended, except for commercial test purposes with the agreement of all

parties involved. If this method is used, it should be stated clearly in

the report.

396
NOTE 5—Load-time curves for tests at 4 K are typically

nonlinear at the start and less regular than load—extension

curves. Also, the effective modulus of a thermally efficient

load train may be low and dependent on the liquid helium level

and the degree of temperature stabilization achieved throughout the system in the temperature range 295 to 4 K. As a re- sult, yield strength data from load—time curves may be less

accurate than those of the recommended method.

8.6.3 Tensile Strength—Calculate the tensile strength by dividing the

maximum load carried by the specimen during the tensile test by the original

cross-sectional area of the specimen.

8.6.4 Elongation—Calculate the percentage increase of elongation

according to Methods E 8 or E 8M, section 7.6.

8.6.5 Reduction of Area—Calculate the percentage of reduction of area

according to Methods E 8 or E 8M, section 7.7.

8.6.6 Rounding Reported Test Data—Round-off the calculated test

results according to Methods E 8 or E 8M, section 7.8.

8.6.7 Replacement Specimens—If necessary, discard data and test

replacement specimens as per Methods E 8 or E 8M, section 7.9.

9. Precision and Bias

9.1 Precision—The precision of these methods is being established.

9.2 Bias—The bias of these methods includes quantitative estimates of

the uncertainties of the dimension-measuring devices, the calibration of

test equipment, and the skill of the operators. At present, bias statement

should be limited to the documented performance or particular laboratories

10. Report

10.1 The test report shall include the following:

10.1.1 Material Characterization—Describe the material tested,

including manufacturing, processing, and metallurgical information.

10.1.2 Specimen Characterization—Describe the specimen location, .m.-i orientation relative to the principal working directions of the stock

report the specimen dimensions, including the cross-sectional size. - let radius, the length of reduced section, and the adjusted length or r:.<- reduced section (if used).

10.1.3 Strain Rate—Report the crosshead speed and the nominal srr.i:::

rate during yielding and after yielding, if a rate change is used.

397
10.1.4

Mechanical Property Measurements—Report the yield strength,

ultimate strength, and method of offset yield strength determination.

Include the method of extension measurement and location of extensometer

attachment, if used. Also report the elongation and its method of calculation, the gage-length-to-diameter ratio for round specimens, and the

reduction of area.

10.2 Optional Data—Report any optional data, such as measurements of

Young's modulus at 4 K, the mechanical properties at room temperature, and

the average grain size of the test material.

10.3 Replicate Tests—If replicate specimens are tested, state the

number of tests, the average values of all mechanical property measurements,

and a measure of scatter.

10.4 Subsize Specimens—If subsize specimens are tested, state any

precautions taken with respect to specimen machining, surface condition, or alignment, and report the grain size of the test material.

10.5 Anomalies—Report any anomalies in material behavior, test

records, or failure type and location.

REFERENCES

(1) Cryogenics Safety Manual

A Guide to Good Practice . 2nd Ed. , Mechanical Engineering Publications, London, 1982.

(2) Basinski, Z. S., "the Instability of Plastic Flow of Metals at Very

Low Temperatures," Proceedings of the Royal Society . Vol. A 240, 1957,

pp. 229-242.

(3) Ogata, T. , Ishikawa, K. , and Nagai, K. , "Effects of Strain Rate on the

Tensile Behavior of Stainless Steels, Copper, and an Aluminum Alloy at

Cryogenic Temperatures," Transactions of the Iron and Steel Institute

of Japan . Vol. 71, No. 10, 1985, pp. 1390-1397. Te fsa - to - H* ,Kr

(4) Ogata, T. and Ishikawa, K. "Time -dependent Deformation of Austenitic

Stainless Steels at Cryogenic Temperatures," Cryogenics . Vol. 26,

1986, pp. 365-369.

(5) Schwartzberg , F. R. , "Mechanical Property Testing Techniques for the

Cryogenic Temperature Range," in: Advances in Cryogenic Engineering . Vol. 8, Plenum Press, New York, 1963, pp. 608—623.

(6) Evans, D. , Simmonds, G. E. , and Stapleton, G. B., "Improved Facility

for Determining Mechanical Properties of Materials in Liquid Helium,"

in: Proceedings. Fourth International Cryogenic Engineering Conference . IPC Science and Technology Press, Guildford, Surrey, England,

1972, pp. 331-337.

398
( 7 ) Hartwig, G. and Wuchner, F. , "Low Temperature Mechanical Testing

Machine," Review of Scientific Instruments . Vol. 46, 1975,

pp. 481—485.

(8) Novikov, N. V. "Mechanical Property Measurement Techniques of Structural Materials at Cryogenic Temperatures," in: Advances in Cryo gen i:

Engineering . Vol. 22, Plenum Press, New York, 1977, pp. 113—118.

(9) Horiuchi, T. , Shimada, M. , Fukutsuka, T. , and Tokuda, S., "Design and

Construction of an Apparatus for Testing Materials at Cryogenic

Temperatures," in: Proceedings. Fifth International Cryogenic

Engineering Conference . IPC Science and Technology Press, Guildford,

Surrey, England, 1977, pp. 465—468.

(10) Sparks, L. L. , "Temperature, Strain, and Magnetic Field Measurements,"

in: Materials at Low Temperatures . R. P. Reed and A. F. Clark, eds .

American Society for Metals, Metals Park, Ohio, 1983.

(11) Read, D. T. and Tobler, R. L. , "Mechanical Property Measurements at

Low Temperatures," in: Advances in Cryogenic Engineering—Material?

Vol. 28, Plenum Press, New York, 1982, pp. 17—28.

399
TABLE 1 Standard Specimens for Room Temperature Tests and

Recommended Proportionally Reduced Standard Small Specimens for 4-K Tests 1

(a)

Metric Versions

G/D ratio - 5 (dimensions, mm)

Standard Specimen

( room- temperature

)

Standard

Small Specimen

(4-K)

Nominal Diameter 12.5 7

G, gage length

D, diameter

R, fillet radius

A, reduced section

62.5 ± 0.1

12.5 ± 0.1

10

75

35 ± 0.1

7 ± 0.1

7 42

(b)

U.S. Customary Versions

G/D ratio - 4

(dimensions, in)

Standard Specimen

( room - temperature

)

Standard

Small Specimen

(4-K)

Nominal Diameter 0.5 0.25

G, gage length

D, diameter

R, fillet radius

A, reduced section

2.000 ± 0.005

0.500 ± 0.010

0.375

2.25

1.000 ± 0.005

0.250 ± 0.005

0.1875

1.25

'See also the notes to Fig. 8 in Methods E 8 and E 8M

.

400
LOAD

Pull Rod

Room Temperature

Load Frame

Vacuum-Insulated

Transfer Tube

Cryogenic

Load Frame

Specimen

a u

U

0

Alignment Device

Electrical

Feed-Through

Vent

Dewar Seal

Vacuum-Insulated

Dewar

FIG 1 Example of a cryostat with Double-Dewar Arrangement

for Tensile Testing at 4 K

401
STRESS,

MPa

~

STRESS,

MPa

C D

FIG 2 Engineering Stress-Strain Curves for an AISI 304L

Austenitic Stainless Steel at 4 K

402

TEMPERATURE,

K

TEMPERATURE,
White _ (- signal) r

Active Gage

120 or 350 a

Strain Gage

(4 K)

Red (+ excitation)

(constant voltage)

4 K

Dummy

Resistor

-VWV\r —

Shunt Resistor

R^

Room Temperature

Dummy Resistor*

Green

(+ signal)

Room Temperature

Dummy Resistor *

Black

(- excitation)

FIG 3 Method of Strain Gage Wiring for Cryogenic Tests

403
STRESS

A

f

G

FIG 4 Round Bar Specimen Configuration (see Table 1)

A B

FIG 5 Stress-Strain Diagram for Determination of Yield Strength

by the Offset Method: A) Serrations occurring after 0.2% strain

B) Serrations occurring before 0.2% strain.

404


U.S. —JAPAN COOPERATIVE PROGRAM COMMITTEE ON TEST METHODOLOGY

FOURTH DRAFT

MARCH 1987

NOTE: This is a working document of a standard being developed by the above

committee. Some technical aspects are still being resolved and are subject to revision. The document in its present state is intended

for committee review purposes, but not for general distribution at

this time.

Proposed Standard Method for

FRACTURE TOUGHNESS TESTING OF STRUCTURAL ALLOYS

AT LIQUID HELIUM TEMPERATURE

1 . Scope

1.1 This standard describes the equipment, procedures, and data analysis used for measuring the fracture toughness of structural alloys in lieu:

helium [4 or 4.2 K (—452°F)]. Procedures described in ASTM Methods E 399 -- and E 813-81 are adopted and modified to address the specific needs and con- cerns of cryogenic testing.

NOTE 1 —The boiling point of normal liquid helium (He I)

at sea level is 4.2 K. This value decreases slightly with

geographic elevation; it is 4.0 K at NBS in Boulder, Colorado,

1677 m (5500 ft) above sea level. In this standard, the test

temperature is nominally designated 4 K.

1.2 To measure the fracture toughness according to this standard,

fatigue-cracked specimen is cooled in a cryostat to 4 K and then load.- i

• a moderate speed while fully submerged in normal liquid helium (He I standard specimen is a pin-loaded, compact specimen, nominally 25 mm i

thick and 50 mm wide.

1.3 This standard specifies the measurement of Kj c or J^ c vhi

appropriate for the load-displacement behavior at 4 K. For hi ;;h- tou. ss alloys, Jj c is determined by a resistance curve technique, but m.eth

measuring resistance cur/e slope, AJ/Aa, are not included in this

1.4 Values stated in SI units a*re treated as primary; values 1 U.S. customary units, secondary.

1.5 This standard does not address safety issues. It is the

responsibility to implement appropriate safety practices Safe* v ;*u i ;•»!: '.«•» for handling liquid helium and other ervogens are available . ;

. . , <• <1

405
2. Applicable Documents

2 .

1

ASTM Standards:

E 4 Practices for Load Verification of Testing Machines

(ANNUAL BOOK OF ASTM STANDARDS, Vols. 03.01, 04.02, 07.01,

and 08.03)

E 8 Methods of Tension Testing of Metallic Materials

(ANNUAL BOOK OF ASTM STANDARDS, Vols. 01.02, 02.01, 02.02,

02.03, and 03.01)

E 8M Methods of Tension Testing of Metallic Materials, Metric

(ANNUAL BOOK OF ASTM STANDARDS, Vol. 03.01)

E 399 Standard Test Method for Plane-Strain Fracture Toughness

of Metallic Materials (ANNUAL BOOK OF ASTM STANDARDS

,

Vol. 03.01)

E 616 Terminology Relating to Fracture Testing

(ANNUAL BOOK OF ASTM STANDARDS, Vol. 03.01)

E 813 Standard Test Method for Jic> A Measure of Fracture

Toughness (ANNUAL BOOK OF ASTM STANDARDS, Vol. 03.01)

3. Definitions

3.1 The terms relating to tensile testing as defined in Definitions

E 616 shall apply in this standard.

3.2 stress intensity factor , plane-strain fracture toughness , and crack

plane orientation—the definitions in Methods E 399 apply in this standard.

3.3 J - integral . physical crack size , original crack site , crack extens ion . effective vield strength , blunting line , initiation of stable crack

growth . regression line slope , and normalized crack size—the definitions in

Methods E 813 shall apply in this standard.

3.4 fracture crvostat—an apparatus consisting of a loading frame and

Dewar for testing specimens in cryogenic environments. A schematic is shown

in Fig . 1

.

3.5 Dewar—a vacuum- insulated container for cryogenic fluids.

3.6 adiabatic heating—the internal heating of a specimen caused by

fracture testing under conditions such that the heat generated by plastic

work cannot be quickly dissipated to the cryogen.

3.7 J c—a critical value of J at the onset of fast fracture causing

large, uncontrolled crack propagation.

406
4 . Summary of Methods

4.1 This standard is based on Methods E 399-83 and E 813-81. From them

procedures are adapted and modified in some respects for use at 4 K. Familiarity with them is assumed.

4.2 The object of this standard is to measure the fracture toughness of

structural alloys at 4 K using a K- or J-based fracture criterion. The type

of load-deflection behavior exhibited by the test material in the thickness

tested determines which fracture criterion is to be used (see Fig. 2).

4.3 The compact specimen configuration is adopted as standard because

it offers a larger fracture toughness measurement capacity for its size than

any of the common fracture specimens. A relatively low load applied to the

compact specimen produces a large stress intensity factor. This enables the mass of the cryostat, grips, and load train to be minimized, lowering the

cryogen consumption and cost per test.

4.4 To determine the fracture toughness by this standard, a fatiguecracked specimen is cooled to 4 K in a suitably designed cryostat containir.;:

liquid helium. Using displacement control, the specimen is loaded while the

load versus displacement curve is autographically recorded. The material

response determines the appropriate method of analysis.

4.4.1 The parameter Kp c is used to characterize materials with a Type T record. The loading is linear-elastic and the crack- tip plastic zone is

small compared with the crack size and specimen dimensions. Except for the

provision of appropriate cryogenic apparatus, the method of Kp c measurement

is similar to that set forth in Methods E 399-83.

4.4.2 The parameters J c is used in this standard to characterize Type

II records where the loading is nonlinear and the crack-tip plastic zone is

sizable compared with the crack size and specimen dimensions. This paramet.-:

is not a material constant, but it may be used to designate the toughness

materials that lack sufficient section thickness or toughness for valii - or Jq c tests. The value of J c is calculated from the area under the load- vs -displacement curve up to the point of fast fracture.

4.4.3 The parameter Jj c is used to characterize materials disp lav : r.,; . Type III record where the loading is nonlinear (elastic-plastic, or full-.- plastic), the crack-tip plastic zone is large compared with the crack « and specimen dimensions, and a satisfactory resistance cur/e can be

for the material. The recommended method for Jp c measurement is the

specimen unloading-compliance test procedure.

5. Significance and Use

5.1 At 4 K,

parameter is the room- temperature

the significance and use of the parameter K * as a same as that discussed in Methods E 399-83, sectir: or tests

407
5.2 The parameter J c , as described in this standard, is not a material

constant because its value depends on specimen configuration and size. The

significance of J c values with respect to material behavior or structural

design is not addressed by this standard. The intent of such a measurement

is determine a value of toughness for materials that do not fit either a Kj c or Jj c characterization in practical section thicknesses.

5.3 At room temperature, the significance and use of the parameter Jj c is discussed in Methods E 813-81, section 4. At 4 K, discontinuous plastic

flow may affect the significance of Jj c as a material property. This subject

is currently being evaluated.

5.4 The load-deflection curves of alloys fracture - tested at 4 K typically exhibit instabilities or serrations. The serrations are free -running

events, that may arise from: a) subcritical crack extensions (pop- ins) owing

to brittle cracking in low- toughness materials or b) ductile flow instabilities due to discontinuous plastic deformation in tough materials (2) . Serrations are always accompanied by temporary temperature rise, as shown for a ductile austenitic steel in Fig. 3. Owing to this phenomenon the J- integral

calculations, resistance curve slope, and tearing modulus parameters at 4 K

lack the significance they have at ambient temperatures, where crack growth

is exclusively a slow, stable tearing process.

5.5 The Jj c values measured for high- toughness alloys may be used to

estimate Kj c values following Methods E 813-81, section 9.4:

Kic(J) - (Jic x E ) 4

Such conversions are useful to compare different materials quantitatively on

the basis of a single parameter. These estimates are denoted by the symbol

Ki c (J). Thus, Kj c values are estimated from small, ductile test specimens

that do not actually exhibit linear-elastic fracture or Type I records in

practical section sizes.

5.6 Owing to adiabatic heating, fracture toughness measurements at 4 K

may be affected by rapid loading (3). Therefore, tests at 4 K are performed

at low loading rates (see section 9.4). The potential effects of heating

must be considered in applications subject to high loading rates.

5.7 The fracture parameters Kj c and Jj c are suitable for characterizing

a wide range of materials at 4 K. Nevertheless, some materials cannot be

tested satisfactorily by the methods of this standard. These materials are:

a. too brittle to be properly fatigue-cracked or too thin

(B < 1.6 mm) for K testing (see Methods E 399-83,

section 1.1),

b. of intermediate toughness such that the loading behavior is not

sufficiently linear for a K]; c test, yet fast fracture prevents

the full development of a resistance curve,

c. too tough for a valid J Ic test (see Methods E 813-81, section 4.3).

408
Retesting with a change in specimen dimensions may be necessary to obtain

valid data for some materials, but there is no assurance that valid results

for a given material will be obtained by the procedures of this standard.

6 . Apparatus

6.1 Test Machines—Tests shall be performed in displacement control

with a screw-driven or servohydraulic tensile machine that meets the re- quirements of Practices E 4 regarding the verification of load accuracy.

The equipment must include a means for autographic recording of the applied

loads and specimen displacements. The machine compliance shall be measured

with the cryostat to be used at 4 K; this may be accomplished by replacing

the specimen with a rigid block.

6.2 Cryogenic System Design—For conventional test machines having a maximum load capacity of 10 tons, it is recommended that the cryogenic

apparatus be designed to accommodate a standard, 25-mm-thick compact specimen (see section 8.1.3). A servohydraulic test machine with an inverted

stage (actuator mounted above the specimen) and a cryogenic loading frame

that is free at the bottom end is also recommended. Such a system is nearl

optimum in regard to operating convenience and cost effectiveness for

testing a wide range of materials at 4 K.

6.3 Cryogenic Construction Materials—Many materials, including most

steels, are brittle at 4 K. All cryostat components must be well designed

and fabricated using cryogenic alloys of suitable strength and toughness.

Austenitic stainless steels are generally recommended. Maraging steels,

nickel-plated to prevent rusting, are useful for specimen clevises. The

alpha—titanium alloys are attractive for certain applications requiring hi.:

strength- to- thermal conductivity ratios, such as pull rods.

6.4 Dimension-Measuring Devices—Micrometers and other devices used

for measuring the dimensions of specimens shall be accurate and precise t at least one-half the smallest unit to which the individual dimension is

required to be measured.

6 . 5 Cryostats and Support Apparatus :

6.5.1. Cryostats—A cryostat is required to retain liquid helium : the duration of testing. Usually, cryostats must be custom-built, but • can be designed for use with commercially available, vacuum- insula:

Dewars. Efficient design involves trade-offs between thermal and - considerations (4). Some practical designs are described in the liter.r . • including turret-disc apparatus for testing several specimens sue after a single cooling (4-7).

6.5.2 Dewars—Stainless steel Dewars are recommended for low maintenance, durability, and safety. Glass Dewars are fragile and ha;:.ir-:. -;s and fiberglass Dewars are relatively expensive. A single helium r.*-var should generally be sufficient for tests of short duration. Dnuii i -• -

i <va

409
arrangements using an outer Dewar of liquid nitrogen to surround the inner

Dewar containing liquid helium are also possible (4—6).

6.5.3

Ancillary Equipment—Dewars and transfer lines for handling the

liquid helium must be vacuum- insulated. Pressurized gases, vacuum pumps,

and liquid nitrogen apparatus are therefore required. After testing, the

helium may be released as a vapor to the atmosphere, recycled, or reliquefied. Recycling or reliquefaction requires substantial investment in purification and cryogenic systems.

6 . 6 Clevis Grins :

6.6.1 Design—Conventional clevis grips conforming to the design in

Fig. 3 of Methods E 813-81 may be used. The loading pin holes are designed

with flats to allow specimen rotation during testing. To minimize friction,

the flats must remain free of plastic indentations. Structural alloys with

adequate 4-K strength, toughness, and fatigue resistance are required.

6.6.2 Construction—The maraging steels (i.e., 250 and 300 grades) have

been used for clevis grips of the conventional design (4-7). If materials

of lower strength are used, such as AISI 304LN, the thickness of the clevis

arms should be enlarged.

6 . 7 Extensometer :

6.7.1 Requirements—An extensometer device is required to measure edge

and load-line displacements for Kj c or J tests, respectively. In addition to

rendering accurate displacement measurements at 4 K, the device must be able

to sustain repeated thermal shock, moisture condensation, and mechanical

fatigue

.

6.7.2 Types—Conventional clip gages, such as those described in Methods E 399 and E 813, may be purchased or built using good cryogenic materials and fabrication techniques (8)

.

The clip gage recommended in Methods

E 399 may be used if the total displacement is expected to be 2.5 mm or less.

For a larger linear working range, enlarged gages are necessary (see Fig. 4

of Methods E 813)

.

6.7.3 Construction—A satisfactory combination of active strain gage

elements, backing materials, and bonding agents must be used for constructing cryogenic extensometers . One common choice is a Ni-Cr-alloy gage with a temperature-compensated active element. Typically, 120- or 350-0 gages are wired in a Wheatstone bridge arrangement. Electrical wires insulated with

polytetraf luoroethylene and soft solders with low tin contents are recom- mended to reduce mechanical failures and shorts at 4 K.

6.7.4 Calibration—The clip gage or other extensometer device must be

calibrated at room temperature and at 4 K. This may be done with a barrelmicrometer calibration device fitted with vertical extension tubes with the

gage mounted at the lower end and immersed in liquid helium for calibration

at 4 K (8). Once the calibration is known and proven linear, accurate, and

reproducible, room- temperature checks may be performed to verify the 4-K

410
calibration indirectly. Periodically, the 4-K calibration must be directly

verified, especially when damage is suspected or repair has been performed.

6.8 Temperature Maintenance and Liquid-Level Indicators—The intended

cryogenic environment is ensured by the fixed boiling point of the liquid

helium bath. A thermocouple on the specimen is not necessary, but a meter

or indicator is required to ensure the proper liquid level. An on—off

carbon-resistor meter, a continuous- level meter with superconducting wires,

or a thermocouple ( Chrome 1 -versus -go Id + 0.07% iron) shall be positioned in

the cryostat above the specimen to ensure that it remains fully submerged at

all times during the test.

7 . Sampling

7.1 Material Condition—Samples for specimen preparation shall be taker,

from the material in its final metallurgical condition. All heat treatments

shall be performed prior to sampling and specimen machining.

7.2 Sample Location—The sample location shall be chosen to be representative of the stock material. Allowance must be made for areas affected

by the cutting process.

7.3 Specimen Orientation—The orientation of the specimen notch plane

and the direction of load application relative to the principal stock working directions must be stated in the notation of Methods E 399.

8. Specimens

8 . 1 Specimen Configuration :

8.1.1 Specimen Types—In principle, any specimen configuration cited in

Methods E 399-83 or 813-81 may be used if the standard's guidelines are observed. Current practice at 4 K is limited, however, to compact and bend

specimens. The compact specimen predominates. Bend specimens require more mass for an equivalent ligament size. In addition, the long span of the •• specimen transverse to the loading axis requires cumbersome fixtures th.it

inconvenient for some cryostats.

8.1.2 Standard 4-K Test Specimen—A standard compact specimen is : for tests at 4 K. Its thickness, B, is 25.4 mm; its width, W, is 50.3 mm The W/B ratio is 2.0, and the planar dimensions are proportion.il : :> k

specimen is generally recommended because it is large enough to provide

results at 4 K for most structural alloys, including the austenitic

steels

8.1.3 Alternative Specimens—Selection of specimen size and sha; <• 1 depend on technical purpose, the load limit of the testing machine, r < need to obtain representative samples of the material being tested if standard specimen is not practical, an alternative geometry or selected, following Methods E 399 and E 813.

411
8 .

2

Specimen Preparation:

8.2.1 Starter Notch—The notch should be machined according to Methods

E 813-81.

8.2.2 Loadline Notch Modification—For Jj c measurements , a loadline

notch modification (see Methods E 813-81) for clip-gage attachment between

razor blades is recommended. Integrally machined knife edges may tend to

increase friction effects. If the loadline displacement is not directly

measured in Jj c tests, then the method of determining the displacement at

the loadline from measurements at the edge location must be described.

NOTE 2—If a loadline modification is used and a Type I load—displacement curve is obtained, then Kj c may be calculated as usual.

8 . 3 Fatigue Precracking :

8.3.1 General Requirements—All specimens shall be carefully fatiguecracked. The load range (AP) during fatigue shall not be less than 90% of

the maximum fatigue load. Also, the maximum stress intensity factor during

the fatigue crack extension shall not exceed 0.005 mm"* times the Young's

modulus, E.

8. 3. 1.1 Fatigue precracking at 4 K. If the specimen exhibits linearelastic behavior, the maximum stress intensity factor during fatigue crack

extension shall not exceed 60% of Kj c

. If the specimen exhibits plastic

behavior, the maximum fatigue load must not exceed 0.4P^, where P* is the

limit load. For the compact specimen:

PL - Bb 2 <7y/(2W + a)

8. 3. 1.2 Fatigue precracking at higher temperatures. Fatigue precracking can be performed at a higher temperature if the procedure does not affect the subsequent fracture toughness measurement at 4 K. At the present

time, however, no clear rules have been established for fatigue precracking

at higher temperatures, such as 295 K. If fatigue precracking is performed

at a higher temperature, the maximum loads should be reduced by the ratio of

the yield strength at the fatigue temperature to that at 4 K. If any other

procedure is used, it must be demonstrated that this does not affect the

toughness test result.

NOTE 3 —Sizeable reductions of the maximum load during

fatigue cracking should be avoided because they may retard

the crack growth rate or increase the degree of crack front

curvature

.

8.3.3 Fatigue Crack Length—The ratio of the fatigue crack length to

specimen width, a/W, shall be in the range from 0.45 to 0.70. The choice

of crack length ratio is influenced by the behavior of the test material.

If a Type - 1 test record is expected, or if Kj c may be directly measured,

the crack length ratio shall be 0.45 to 0.55 to maintain compatibility with

412
Methods E 399-83. If nonlinear loading behavior is expected and a resistance curve can be determined, the ratio should be about 0.6; for the

exceptionally strong and tough steels, a ratio near 0.7 may be necessary

to keep within machine load limitations. For alloys of unknown toughness,

a ratio of 0.6 is recommended for initial tests; this can be adjusted for

retests, if necessary.

8 . 4 Side Grooves :

8.4.1 Acceptable Uses—Side grooves should not be used for materials

having Type I or Type II test records. For Type III records, side grooves

may be used to reduce the fatigue crack front curvature, to alleviate crack

front tunneling, or to improve the agreement between the optically measured

and compliance-calculated end-of-test Aa values.

8.4.2 Machining—Side grooves, if used, shall be machined after precracking, following Methods E 813-81, section 7.5.1. For standard 25-mmthick compact specimens, the total depth of side grooving shall amount to

not more than 20% of the gross specimen thickness.

9. General Procedures

9.1 Specimen Measurement—Measure the specimen dimensions according to

Methods E 399-83 (section 8.3) and E 813-81 (section 8.8).

9.2 Installation—Install the specimen in the apparatus, leaving enough

slack for the instrumentation wires so that they will be neither stretched

nor crimped during positioning of the Dewar and testing. Prior to cooling,

check the seating of the specimen and grips at room temperature. A slight

preloading of the specimen is permitted for this purpose.

9.3 Cooling—Prior to cooling, dry the apparatus thoroughly with pressurized air or a heating gun. Also, purge any moisture from the cryogenic

transfer lines. Ice can block the transfer lines or cause erratic load in,;

behavior if it forms on the loading pins, clevises, specimens, clip gage,

load-train members. Next, position the Dewar and precool the apparatus * 77 K by transferring liquid nitrogen into the cryostat. After attain::.,-,

equilibrium at 77 K, remove all the liquid nitrogen and transfer 1: p.:i

helium into the cryostat. Add sufficient liquid helium to compensate : ,r evaporation during the time required to complete the test so that tin- specimen remains fully submerged at all times.

NOTE 4—Compared with liquid helium, the heat transfer

characteristics of gaseous helium are inferior; therefor--,

liquid helium is required to minimize specimen heating.

9 .

Loading :

9.4.1 Control Mode—The test shall be conducted using d 1 :,p 1 a<- trol, not load control. Crosshead displacement is typical for screw -an

413
machines. For servohydraulic machines, stroke control is often used, but

clip-gage control may be a useful alternative for Jj c tests.

9.4.2

Load Rate—Load the specimen at a rate such that the time, tp , to

reach the load at first deviation from the linear portion of the test record

is between 0.5 and 10 min. For other details of test procedure, see Methods

E 813-81, section 8. For standard 25-mm- thick compact specimens, the cross- head rate should be 0.5 mm -min-1 or less.

9.5 Test Record—Autographically record the outputs of load-cell versus displacement following the guidelines of Methods E 399-83, section 8.4

for Kj c tests, or as appropriate for Jj c tests. For Jj c tests, a real-time

computer-calculated J—Aa curve record is also recommended.

10. Interpretation and Calculation of Results

10.1 General—Classify the test record as one of the three basic types

shown in Fig. 2. Depending on the load-versus-displacement behavior, decide

which fracture criterion is to be used. More than one fracture parameter

may be applied in some cases.

10.2 measurement—For Type I test records showing linear-elastic

behavior, Kj c should be calculated according to Methods E 399. Construct

the secant offset line, calculate the conditional toughness value, Kq , and

determine its validity as a Kj c value according to Methods E 399-83, section

9. If the test fails to qualify as a valid Kj c measurement, calculate the

value of J c according to the following section, 10.3.

10.3

J c measurement—For Type II tests that produce nonlinear loadversus -deflection curves, determine a value of J c as follows: Calculate the

total value of J corresponding to the total area under the load-displacement

curve up to the point of fast fracture. The area to be measured is cross- hatched in Fig. 2. The measurement shall be accurate to 2% and it may be

performed using a polar planimeter or numerical integration of the loaddisplacement signals by computer techniques. Calculate the value of Jc for

compact specimens:

NOTE 5—Notch cut-out modifications (instead of the usual

edge-displacement measurements) do not violate the Kj c calculations at 4 K, according to this standard.

J c - f (a 0 /W) -A/Bb

where

f(a 0 /W)

original crack size, including precrack,

specimen width,

a dimensionless coefficient that corrects for the tensile

component of loading, as given in section 9.1.1 and Table 1

414
A => area under the load versus load-point displacement record

in energy units,

B = specimen thickness, and

b = initial uncracked ligament, W — a 0 .

10.4 J Ic measurement—For Type III tests, Jj c should be measured by a resistance curve technique, where J is plotted versus physical crack growth,

Aap . The single-specimen unloading compliance method is recommended as the

most suitable technique at 4 K because it reduces the number of specimens

required and the cryogenic fluid costs. Guidelines for the method are given

in the appendix of Method E 813-81.

10.4.1 Calculations of the J- integral are made from the load vs loadpoint displacement curves by using the expression in section 10.2. For side

grooved- specimens , the net thickness is used in the J calculations and effective thickness is used for crack growth predictions from unloading com- pliance as per Methods E 813-81 sections 8.7.1 and 9.1.1.

10.4.2 Prepare a plot of J versus Aa

p spanning the 0.15—1.5 mm exclusion interval, as given by Methods E 813-81. Calculate the area under the

load—displacement curve, including the area under serrations. For tests in

clip-gage displacement control, the area under a serration is zero. Thus,

tests in stroke control are expected to give higher values of J for a given

crack extension, owing to the positive increase of area associated with ser- rations. The intent is not necessarily to measure a physically meaningful

J-Aa curve, but just a regular data set enabling linear extrapolation to the

initiation of crack extension at Jj c

.

10.4.3 Construct the blunting line and plot the regression line, and

qualify the J—Aa calculations according to Methods E 813-81, section 9.2.

10.4.4 Determine Jq as the value of J at the blunting and regression

line intersection. Qualify the data and determine whether Jq value is a valid Jj c determination according to Methods E 813-81, sections 9.3 and A'.

11. Report

11.1 Material Characterization :

11.1.1 Material Description—Describe the material tested. incA;d;:

the product form (e.g., plate, forging, casting), processing and heat tnu:

ment, and any other relevant metallurgical information.

11.1.2 Tensile Properties—Report the yield strength, tensile

and elastic modulus at room temperature and at 4 K.

11.2 Specimen Characterization :

11.2.1 Describe the crack-plane orientation according to *':.<• M< •. E 399-83 identification code.

415
11.2.2 Report the specimen dimensions, including thickness, width, and

ligament and length crack lengths.

11.2.3 If side grooves are used, describe the notch geometry (included

angle, notch-tip radius, and net section thickness).

11.3 Fatigue Cracking Characterization :

11.3.1 Report the maximum fatigue load, the maximum stress - intensity

factor, the stress - intensity factor range, and the temperature used for the

fatigue precracking.

11.3.2 Report the final fatigue crack length obtained from the three - point average for Kj c tests or the nine-point average for Jj c tests, as described in Methods E 399-83 or E 813-81. Also measure the surface crack

lengths. Report the maximum deviation of any single measurement from the

average

.

11.4 Test Record—Describe the load-versus -displacement record or designate the type of test record observed according to the examples shown in

Fig. 2.

11.5 Loading Rate—Report the loading rate in terms of tp or the rate

of increase of the stress - intensity factor, K, during the linear region of

loading. Specify whether crosshead speed, stroke displacement, or clip-gage

displacement was used.

11.6 Fracture Toughness Measurements :

11.6.1 For a Type I test record, report the fracture toughness in terms

of Kq or Kj c and include an assessment of the validity according to Methods

E 399-83.

11.6.2 For Type II test records, where a full resistance curve cannot

be developed, report the value of J c as defined in section 3 of this standard. If an unloading compliance technique was used for J—Aa measurements

up to the point of fast fracture, estimate the value of crack extension at

the J c measurement point.

11.6.3 For Type III test records, where a full resistance curve can

be developed, report the value of Jj c and calculate Kj c (J). Assess the

validity of Jj c according to the criteria of Methods E 813-81.

11.7 Fracture Surface Appearance—Describe the fracture appearance.

For computerized Jj c tests, report the end-of-test crack extension increment

and the accuracy of the compliance-predicted versus the optically measured

final average crack length.

11.8 Replicate Tests—Give the number of replicate tests performed, the

average fracture toughness values, and a measure of scatter.

416
11.9 Anomalies—Report any anomalies observed in material behavior, the

test records, or failure mechanisms.

12. Accuracy and Precision

12.1 Accuracy—There is no accepted standard fracture toughness value

for any material. In the absence of such a true value, any statements about

accuracy are meaningless.

12.2 Precision—Information concerning the precision of 4-K fracture

tests is being obtained from interlaboratory test programs in progress.

REFERENCES

(1) Cryogenics Safety Manual—A Guide to Good Practice . 2nd Ed., Mechanical Engineering Publications, London, 1982.

(2) Tobler, R. L. and Shepic, J. A. "Design and Performance of a Ringshaped Clip Gage for Fracture Mechanics Testing," Journal of Testing

and Evaluation . Vol. 13, 1985, pp . 299—302.

(3) Tobler, R. L. , Ohnishi , K. , Shoj i , T., and Takahashi, H. , "Fracture,

Acoustic Emission, and Adiabatic Heating of Austenitic Stainless

Steels at Liquid Helium Temperature," in: Progress in Acoustic

Emission III . Japanese Society for Non-Destructive Inspection, Tokyo.

1986, pp. 453-461.

(4) Fowlkes, C. F. and Tobler, R. L. , "Fracture Testing and Results for

a Ti—6A1—4V Alloy at Liquid Helium Temperature," Engineering Fract-:r-~- Mechanics . Vol. 8, 1976, pp . 487—500.

(5) Read, D. T. and Tobler, R. L. , "Mechanical Property Measurements it Low Temperatures," in: Advances in Cryogenic Engineering—' *

. • • •

’ Vol. 28, Plenum Press, New York, 1982, pp . 17—28.

(6) Juhas, M. C., Duncan, D. B., and Streit, R. D., "Development of a Facility for Fatigue and Fracture Testing of Stainless Steels .it Liquid Helium Temperature," Cryogenics . Vol. 23, 1983, pp . 603-';

(7) Shimada, M. , Ogawa, R. , Moriyama, T., and Horiuchi, T., "Deve Lopm^:

of a Cryogenic Fracture Toughness Test System," Crvo-.y-r. : - Vol. 21, No. 5, 1986, pp . 269-274.

(8) Tobler, R. L. , "Ductile Fracture with Serrations in AIST 11 S Stainless Steel at liquid Helium Temperature," in: E 1 a s - j - p . • Fracture: Second Symposium . Vol, II —Fracture

Engineering Applications . ASTM STP 803, C. F. Shih and V eds., American Society for Testing and Materials, Phi lade 1 ;-.h .

- pp. 763-776.

417
(9)

Shimamoto, S., Nakajima, H. , Yoshida, K. , and Tada, E. , "Requirements

for Cryogenic Structural Alloys for Superconducting Magnet Cases," in:

Advances in Cryogenic Engineerine—Materials . Vol. 32, Plenum Press,

New York, 1986, pp . 23—32.

(10) Clarke, G. A., Andrews, W. R. , Paris, P. C., and Schmidt, D. W.

,

"Single Specimen Tests for Jj c Determination," in: Mechanics of Crack

Growth, ASTM STP 590, American Society for Testing and Materials,

Philadelphia, 1976, pp . 27—42.

(11) Joyce, J. A. and Gudas , J. P., "Computer Intractive Jj c Testing of

Navy Alloys," in: Elastic—Plastic Fracture . ASTM STP 668, J. D.

Landes, J. A. Begley, and G. A. Clarke, eds . , American Society for

Testing and Materials, Philadlephia , 1976, pp . 451—468.

(12) G. A. Clarke, "Single-Specimen Tests for Jj c Determination

Revisited," in: Fracture Mechanics: Thirteenth Conference . ASTM

STP 743, R. Roberts, ed. , American Society for Testing and Materials,

1983, pp. 553-575.

418
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

FIG 1 Example of a Cryostat with Double-Dewar Arrar.

for Tensile Testing at 4 K

419
I. Linear-Elastic Behavior: K,c

is measured

II. Elastic-Plastic Behavior: J is measured

III. Fully Plastic Behavior: J|C

is measured

(sizable plastic zone)

FIG 2 Load—Displacement Behavior Observed in Fracture-Toughness Tests

420
Fe —18 Cr — 16 N i -6.5Mn -2.4 Mo

Weld Deposit

T = 4K

Deflection, d, mm

FIG 3 Typical Compact Specimen Loading Behavior at 4 K, Sh v Serrations and Points of Unloading Compliance Method

421



FRACTURE, ACOUSTIC EMISSION, AND ADIABATIC HEATING

OF AUSTENITIC STAINLESS STEELS AT LIQUID HELIUM TEMPERATURE*

R. L. Tobler,t T. Shoj i , H. Takahashi, and K. Ohnishi

Research Institute for Strength and Fracture of Materials

Tohoku University

Sendai, Japan

Conventional fracture mechanics tests supplemented with

acoustic emission and adiabatic heating measurements were con- ducted to characterize commercial austenitic stainless steels at

4 K. Three alloys, representing low, medium, and high toughness

alloys, were examined, and high levels of acoustic emission are reported for each one. The high acoustic emission of these aus- tenitic steels at 4 K is attributed to low temperature strengthening effects in general and to the peculiar failure mechanism

involving discontinuous deformation and adiabatic heating that

occurs in ductile materials near absolute zero.

INTRODUCTION

Detailed fracture characterization of austenitic stainless steels it

extreme cryogenic temperatures is necessary for alloy development and . • conducting magnet design in fusion energy research. Fracture toughness

at 4 K are usually obtained by critical stress intensity (Kj c ) or

(Jj c ) measurements performed with conventional 25-mm- thick compact s-. submerged in liquid helium. Such data describe macroscopic fracture or. giving no information regarding the thermodynamic or micromechan i r

.

'. of fracture processes at 4 K.

Acoustic emission has been previously used to monitor the ope r a *

; r. fully assembled superconducting magnet systems at cryogenic temper a and to study the microdeformation and fracture behavior of f ibe r - r<- : : :

— Published in Acoustic Emission III . The Japanese Societv of .. :r

• Inspection, Tokyo, 1986; pp . 453—461.

(Visiting fellow from the Fracture and Deformation Division, N'.it : of Standards, Boulder, Colorado, USA.

423
plastic composites at temperatures from 295 to 4 K. 3 The purpose of this

study is to explore acoustic emission and adiabatic heating monitoring as techniques to evaluate the 4-K fracture mechanisms of austenitic stainless

steels in greater detail. The fact that abrupt discontinuous yielding with

audible sound is common in these materials at cryogenic temperatures suggests that sensitive acoustic emission techniques may be especially appropriate for such purposes.

EXPERIMENTAL PROCEDURE

Three fully austenitic stainless steels were tested in this study,

including two wrought base metals and a shielded metal arc (SMA) butt weld.

This selection of materials enabled a wide variety of load—displacement

(P—5) behaviors to be evaluated. The wrought base metals were 25-mra- thick

plates of Fe—18Cr—3Ni—13Mn and AISI 310S base metal (BM) tested in the

annealed condition. The SMA weld, designated AISI 310S WM, was prepared

commercially using a type 310S electrode with two 22-mm-thick base plates

and tested in the as -welded condition. The chemical compositions and

tensile properties are listed in Tables 1 and 2.

Compact specimens (nominally 1TCT size) were prepared with notches

machined in the longitudinal plate or weld bead direction (TL orientation)

.

The specimens were 50 mm wide and 22 or 25 mm thick. The Fe—18Cr—3Ni— 1 3Mn

specimens were precracked to a relative crack length (a/W) of 0.5, but they

were not side-grooved in view of linear elastic test requirements. The AISI

310S weld- and base-metal specimens were precracked to a/W - 0.6 and were

side-grooved 20% for elastic—plastic and fully plastic tests.

The specimens were fatigue -cracked at 295 K, then side-grooved or instrumented, as required. Subsequently, they were cooled in a liquid helium

cryostat and fractured using a 10- ton screw-driven machine 4,5 at crosshead

speeds of 1.0 or 0.1 mm/min. After testing, the fracture surfaces were

examined by scanning electron microscopy (SEM) to verify the failure mechanisms and to determine the average dimple diameter and the size of the

stretch zone between the fatigue crack and fracture zones. The specimen

strength ratio was calculated as defined in the ASTM E 399-85 test method.

The tearing modulus, Tm , and plane strain toughness values, Kj c and Jj c , are quoted in the text without further elaboration, since the details are published elsewhere. 5

' 7

An existing acoustic emission system 8 was modified and used for tests

at 4 K by employing a waveguide and reducing the gain. The waveguide was a 4-mm-diameter stainless steel rod, welded to the specimen back face as shown

in Fig. 1. The guide transmitted acoustic emission signals from the specimen at 4 K to a wideband piezoelectric transducer at 295 K via a coupling

outside the cryostat. A significant problem at 4 K is signal saturation

owing to the high energy levels of emissions from these steels. Therefore

our tests used a 20 dB gain, reduced from the usual 60 dB gain at 295 K

The signals were fed into a distribution analyzer, which counted pulses

according to their peak amplitudes and determined the individual event

energies, EAE .

424
Adiabatic heating was detected by measuring the internal temperature,

T, at the crack tip at the specimen midthickness using Chromel-vs -gold-plus - 0 . 07%- iron-alloy thermocouples embedded in the ligament along the fracture

plane. Typically, four thermocouples (0.24-mm wires) were inserted through

2-mm holes drilled 3 mm apart, as shown in Fig. 1. These holes were then

filled with fine Cu powder. The four thermocouple outputs were recorded on

a strip chart, and as cracking propagated across the ligament, the highest T or AT (T — 4.2 K) provided a measure of heating.

RESULTS

The Fe—1 8Cr—3Ni—1 3Mn , AISI 310S WM, and AISI 310S BM steels of this

study displayed typical linear elastic, elastic—plastic , and fully plastic

fracture behaviors with Kj c values of 71, 159, and 193 MPa-m^, respectively.

Scanning-electron-microscope views of the fracture surfaces after testing

are shown in Fig. 2, and typical P—5, acoustic emission, and adiabatic heating behaviors at a crosshead speed of 1.0 mm/min are shown in Fig. 3. The

results, including fracture toughness properties, loading parameters, and

failure mechanisms, are summarized in Table 3.

Linear Elastic Fracture

The Fe—18Cr—3Ni—13Mn steel at 4 K demonstrates brittle fracture behavior. The ratio of the net section fracture stress to the yield strength.

RgC, is 0.5. The failure mechanisms was transgranular cleavage on (111)

planes, as shown in Fig. 2A and previously confirmed by X-ray diffraction

~

The P—5 curves are only slightly nonlinear prior to complete fracture

at the maximum load, Pmax . A high level of acoustic emission is observed

beginning at low loads and continuing up to pmax ( Fi S- 3A) . On the other

hand, adiabatic heating is zero until the terminal stages of loading; then

a small temperature rise, to 7 K, occurs owing to a minor pop -in, followed

by a more substantial rise to 22.5 K during the final fracture (Fig. 33

Elastic—Plastic and Fully Plastic Fracture

Both the AISI 310S weld- and base-metal steels at 4 K display the

essential aspects of ductile fracture; highly nonlinear P—5 curves. Rgratios greater than unity, and dimpled fracture surfaces corresponding, t< microvoid coalescence (MVC) processes (Fig. 2B,C). Compared with the w«- 1

•: metal, the base metal has significantly higher toughness and smaller, r

• numerous serrations. The fully plastic behavior (base metal) is i :

• guished from elastic—plastic behavior (weld metal) by its approach to 1 load failure and its higher resistance to instability. The base ti-r.i. , shows a higher displacement at maximum load, a higher Tm , a larger • zone size, and a larger average dimple diameter (Table 3).

In this paper, ductile fracture is viewed as follows: First,

blunting occurs via stretch zone formation with microvoid or mi r - : > r initiation. Ductile tearing then ensues as a macrocrack forms .mb .m -.,:

425
by linking with growing voids. At 4 K, this process is accompanied by discontinuous loading, which gives the P—5 curves a serrated appearance. Each

load drop is a free-running event that occurs in milliseconds, accompanied

by the evolution of heat and acoustic emission with audible sound. The load

drops are faint at first but steadily increase as testing proceeds.

Continuous loading: During the initial stages of testing, the behavior

featured stable deformation with smoothly rising loads. As nonlinear loading developed, the acoustic emission activity rose, owing presumably to

plastic zone growth, microfracture events, or both. Adiabatic heating also

developed during nonlinear loading, but only at a crosshead speed above some critical value. At the 1.0 mm/min speed, the maximum temperature during

AISI 310S weld- and base-metal tests reached 4.8 and 8.0 K, respectively.

On the other hand, at the lower speed of 0.1 mm/min (data not shown), the

cooling from the cryogen was sufficient to prevent adiabatic heating during

continuous loading.

As shown in Fig. 4A, it is possible to correlate the mutual rise of

acoustic emission and temperature in terms of crosshead speed. In this

graph, the change of the total acoustic emission energy per unit time

(AEE^g/At) is plotted versus AT/At, with data for both crosshead speeds.

This result indicates the existence of a threshold acoustic emission rate

at about 8.5 x 10" 2 V

2 s/s, below which adiabatic heating during continuous

loading was zero for the AISI 310S steels of this study.

Discontinuous loading: As testing continued, load drops formed with

increasing magnitude, coincident with high level bursts of acoustic emission

and adiabatic heating. The temperature rise associated with each serration

was transient and quickly fell to background levels after the discontinuous

deformation was arrested. The temperature rise increases with AP, independent of material type for AISI 310S weld and base metal.

The acoustic emission bursts during serrations usually exceeded the

signal saturation limits, even for a 20 dB gain. For the case of nonsaturated signals in an AISI 310S weld-metal test, however, a correlation between acoustic emission energy and AT for individual events is demonstrated

in Fig. 4B.

DISCUSSION

On the basis of previous experience in general, 9 * 10 annealed austenitic

stainless steels at 295 K must be classified as poor acoustic emission signal emitters. At 4 K, however, the situation changed favorably: all alloys

were significantly strengthened, their ratios of tensile yield strength to

ultimate tensile strength increased, and the process of serrated yielding

with adiabatic heating was prominent in the e lastic—plastic and fully plastic alloys. Reasoning inductively from the three examples of this study, we predict that most austenitic stainless steels will be acoustically active at

4 K. The acoustic emission technique is, therefore, an available tool for

microfracture studies in these steels at 4 K.

426
The identification of acoustic emission sources by metallurgical

methods was not pursued in this study, but the existence of correlations

between acoustic emission and adiabatic heating implies that these parameters describe the same physical processes. In any case, the diverse fail- ure mechanisms of these alloys provide a key to understanding the adiabatic

heating results. In the high strength Fe—18Cr—3Ni—13Mn steel, cleavage

fracture at 4 K was the result of an exceptionally high nitrogen content

(0.37 wt.% N) . In the medium strength AISI 310S weld- and base-metal materials, dimpled fracture occurred via the growth and coalescence of voids

initiated at inclusion—matrix interfaces. The weld metal toughness was compromised, however, by a high inclusion content derived from welding,

and thus, many more voids were initiated in the weld metal. The distance

between voids was relatively small, and, compared with the base metal, less

plastic work was required to cause fracture. Thus, the failure mechanisms

determine the levels of toughness.

The maximum temperatures for the three alloys during their instability

events were 22.5, 53, and 30.2 K (see Fig. 3). Since the corresponding

crack extension increments were dissimilar, comparison is best made in terms

of the normalizing parameter T/Aa. The three Aa increments are 24, 5.5, and

0.08 mm, respectively, where the first value is the ligament length for Fe—

18Cr—3Ni—13Mn steel, and the other two values for AISI 310S weld- and basemetal steels were estimated from compliance changes made before and after

serrations. On this basis, the adiabatic heating behavior during discontinuous loading is compared in Fig. 5. Also shown are the data for continuous loading. We suggest that adiabatic heating should correlate with the

level of plastic deformation associated with fracture, as reflected bv vari- ous physical and mechanical parameters. In Fig. 5, stretch zone size and

are taken to characterize crack initiation resistance; dimple size and tearing modulus function similarly for crack propagation.

Clearly, continuous loading contributes only a modest level of hear in,-

and that can be totally eliminated by the appropriate selection of crossht-.i

speed. But the huge temperature spikes that develop during serrations in

plastic alloys are independent of crosshead speed at normal rates; they cn: not be avoided, and their effects on fracture may be significant. The r <

• vance of this finding to standard 4-K test procedures for elastic-p 1 a ;

• and fully plastic materials must be considered in future work.

CONCLUSIONS

High levels of acoustic emission are emitted during fracture

of austenitic stainless steels at 4 K. Therefore, acoustic emission • niques appear to be useful for microfracture evaluations of these mare:

In this study, two distinctly different correlations between .icons: i

-

• - sion and adiabatic heating were demonstrated for the continuous and : continuous deformation stages of ductile fracture. Significant ad:..

heating occurred, and internal temperature rises of up to 5 3 K re- ported. Apparently, the magnitude of heating can be explained in

the level of plastic deformation associated with fracture427
ACKNOWLEDGMENTS

This work was sponsored by the Ministry of Education of Japan

(Monbusho) , with the first author being supported by the U.S. Department

of Energy, Office of Magnetic Fusion Energy. The authors are also indebted

to Mr. K. Nakano (Tohoku University) for performing some of the tests.

REFERENCES

1. R. S. Kensley, K. Yoshida, H. Tsuji, and S. Shimamoto, Cryogenics, 23

(1983) 17-25.

2. J. Lore, N. Tamada, 0. Tsukamoto, and Y. Iwasa, Cryogenics 24 (1984)

201-207.

3. T. Nishiura, K. Katagiri, S. Owaki, and T. Okada, Cryogenics, 24 (1984)

329-33.

4. M. Takeuchi, T. Shoj i , H. Takahashi, and T. Anayama , Reports of the

Research Institute for Strength and Fracture of Metals, Tohoku

University, Sendai, Japan, 17 (1984) 1—12.

5. M. Takeuchi, T. Shoji, H. Takahashi, and T. Anayama, Trans. JSME (A),

51 (1985) 2256-2264 (in Japanese).

6. K. Ohnishi, T. Shoji, and H. Takahashi, Trans. JSME (A), 52 (1986)

170—173 (in Japanese).

7. R. L. Tobler and D. Meyn, in: "Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures—VIII," NBSIR 85-3025, National

Bureau of Standards, Boulder, Colorado (1985) 167—177.

8. H. Niitsuma, M. Kikuchi, H. Takahashi, M. Suzuki, and R. Sato, in:

"Proceedings of the Fifth International Acoustic Emission Symposium,"

Japanese Society for Non-destructive Inspection, Tokyo (1980) 411—420.

9. T. Ingham, A. L. Stott, and A. Cowan, Int. J. Press. Ves. Piping 3 (1975) 247-293.

10.

H. N. G. Wadley, C. B. Scruby, and J. H. Speake, Int. Met. Rev., Review

249, No. 2 (1980) 41-64.

428
Table 1. Chemical Compositions of Test Materials (wt.%)

Material Cr Ni Mn C Si P S N Mo

Fe—18Cr—3Ni—13Mn 18.09 3.26 13.32 0.038 0.52 0.028 0.005 0.37 0.12

AISI 310S WM 24.68 20.96 2.10 0.10 0.48 0.002 0.003 — —

AISI 310S BM 24.58 19.64 1.52 0.06 0.37 0.013 0.003 — —

Table 2. Tensile Properties of Test Materials

Material

Temperature

(K)

Yield

Strength

Cy (MPa)

Ultimate

Tensile

Strength

ct u (MPa)

°y/aw

Elongation

(%)

Fe—18Cr—3Ni—13Mn 295 440 796 0.55 56

4 1540 1811 0.85 4

AISI 310S WM 293 319 577 0.55 26

4 819 1048 0.78 12

AISI 310S BM 293 225 558 0.40 80

4 796 1322 0.60 52

Table 3. Load Curve Characteristics, Fracture Properties,

and Failure Mechanism Parameters at 4 K

Material

Curve

Type

Load

Drops

ay/au RSC

J Ic (kJ/m2

)

K Ic (MPa-m4

) Tm

Failure

Mechanism

Zone

Size

(pa)

Fe—18Cr—3Ni—13Mn LE one 0.85 0.5 24 71 0 cleavage 0 AISI 310S WM EP several 0.78 2.2 118 159 23 MVC 10 AISI 310S BM FP many 0.60 2.0 333 293 66 MVC 100

LE - linear elastic; EP - elastic-plastic; FP - fully plastic; MVC - microvoid c na . o st r

429
Fig. 1. Nominal 1TCT specimen with instrumentation

(dimensions in millimeters).

430
55 Mm 20 um 20 pm

Fig. 2. SEM fractographs of the

A, Fe—1 8Cr—3Ni—1 3Mn ; B,

4-K fracture surfaces:

AISI 310S WM; C, AISI 310S BM

431
Figure 2. Load-versus-displacement records for three steels tested in liquid

helium at 4 K, showing AE data and internal temperature response.

(AE gain, 20 dB; crosshead speed, 1.0 mm/min)

.

432

AE

EVENT

ENERGY.
Temperature Rise Rate. AT/j.t, K/s

A

Fig. 4. Acoustic emission and adi.iba;

heating correlations

(A) for continuous loading.

(B) for serrations.

B

433
20

Continuous Loading

Fig. 5. Adiabatic heating dependence on toughness parameters.

(A) maximum temperature observed in continuous loading;

(B) maximum temperature per unit crack extension during

final fracture or serration events.

0: Fe—18Cr—3Ni—1 3Mn; A: AISI 310S WM; : AISI 310S BM

434
NBS-114A (REV. 2 -80 ) U.S. DEPT. OF COMM. 1. PUBLICATION OR 2. Performing Organ. Report No. 3. Publication Date BIBLIOGRAPHIC DATA

SHEET (See instructions)

REPORT NO.

NBSIR 87-3067 May 1987

4. TITLE AND SUBTITLE

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures

X

5. AUTHOR(S)

R.P. Reed, Editor

6. PERFORMING ORGANIZATION (If joint or other than NBS, see in struction s) 7. Contract/Grant No.

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE 8. Type of Report Perioo C WASHINGTON, D.C. 20234

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street, City. State. ZIP)

Department of Energy

Office of Fusion Energy

Washington, DC 20545

10. SUPPLEMENTARY NOTES

Previous reports in this series: NBSIR 86-3050, NBSIR 85-3025, NBSIR 84-3000,

NBSIR 83-1690, NBSIR 82-1667, NBSIR 81-1645,

NBSIR 80-1627, NBSIR 79-1609, NBSIR 78-884

[~~| Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a signr can:

bi bliography or literature survey, mention it here)

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

properties of materials that may be used in cryogenic structures for the superconducting magnets of magnetic fusion energy power plants and prototypes. Its

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

jointly by the staffs of the National Bureau of Standards and the Office of Fusi n Energy of the Department of Energy; it is managed by NBS and sponsored by DOE.

Research is conducted at NBS and at other laboratories through subcontracts

with NBS.

Research results for 1986 are presented in technical papers under fiveheadings that reflect the main program areas: Structural Alloys, Weldinp, N

• - metallics, Technology Transfer, and United States-Japan Cooperative Prcprin.

Objectives and research highlights are summarized in the introduction to ea< program area.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and ;ep<jrctr

austenitic steels; composites; cryogenic properties; elastic properties; v i- tures; mechanical properties; nonmetallics ; stainless steels; s t ru<- 1 -i r.i

13. AVAILABILITY 7; "7^ RlN PA QF j |X^ Unlimited

| | For Official Distribution. Do Not Release to NTIS

I I Order From Superintendent of Documents, U.S. Government Printing Office. <Vash rngtor. . z.' . » — 20402.

[XX Order From National Technical Information Service (NTIS), Springfield, VA, 22>b