MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES - VI

To

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

By

Fracture and Deformation Division

National Bureau of Standards

Boulder, CO 80303

—QC —

100

. U56

33-1690

1983

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NBSIR 83-1690

V 4. % TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES

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Edited By

R.P. Reed and N.J. Simon

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

May 1983

Sponsored By

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

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

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

CONTENTS

£iSi

SUMMARY v

ORGANIZATIONAL CONTACTS vi

PROGRAM DESCRIPTION vi

i

HIGHLIGHTS OF RESULTS 1

STRUCTURAL ALLOYS

STRUCTURAL ALLOYS PROGRAM 9

SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS

OF STRUCTURAL ALLOYS 10

MARTENSITIC TRANSFORMATIONS IN Fe-Cr-Ni STAINLESS STEELS 11

A COMPUTER- INTERACTIVE FATIGUE CRACK GROWTH RATE TEST

PROCEDURE 41

EFFECTS OF SENSITIZATION ON SOME FRACTURE MECHANICS FATIGUE

PARAMETERS FOR Fe-18Cr-3Ni-13Mn-0.37N STAINLESS STEEL AT 295

AND 4 K 53

EFFECTS OF MAGNETIC FIELD ON TENSILE BEHAVIOR AT 4 K OF ALLOYS

304 AND 310 73

FACTORS INFLUENCING THE LOW-TEMPERATURE DEPENDENCE OF YIELDING

IN AISI 316 STAINLESS STEELS 87

TEMPERATURE DEPENDENCE OF FLOW STRENGTH -OF SELECTED AUSTENITIC

STAINLESS STEELS 113

PROPERTIES OF AUSTENITIC STAINLESS STEEL AT CRYOGENIC

TEMPERATURES 127

AUSTENITIC-STEEL ELASTIC CONSTANTS 157

ELASTIC CONSTANTS OF AN Fe-5Cr-26Mn AUSTENITIC STEEL, 76-400 K. . 181

WELDMENTS AND CASTINGS

WELDMENTS AND CASTINGS PROGRAM 197

CORRELATIONS OF FATIGUE CRACK GROWTH RATE PARAMETERS AT

CRYOGENIC TEMPERATURES 199
CONTENTS (Cont'd)

Page

DEFORMATION AND FRACTURE OF STAINLESS STEEL CASTINGS AND

WELDMENTS AT 4 K 213

MECHANICAL PROPERTIES OF STAINLESS STEEL CASTINGS AT 4 K 235

METALLOGRAPHY OF DUPLEX STAINLESS STEEL CASTINGS 243

A TRANSMISSION ELECTRON MICROSCOPY STUDY OF CAST DUPLEX

AUSTENITIC STAINLESS STEELS WITH VARYING FERRITE CONTENT:

AS-CAST AND DEFORMED AT 4 K 273

TEXTURE IN STAINLESS-STEEL WELDS: AN ULTRASONIC STUDY 339

NONMETALLICS

NONMETALLICS PROGRAM 359

VARIABILITY IN MECHANICAL PERFORMANCE OF G-IOCR CRYOGENICGRADE INSULATING LAMINATES 361

EFFECT OF CRYOGENIC TEMPERATURES ON THE MECHANICAL PERFORMANCE

OF GLASS-FABRIC-REINFORCED EPOXY AND POLYMIDE MATRIX LAMINATES. . 369

ELASTIC STIFFNESS AND LOCAL STRESSES IN WOVEN-FABRIC

COMPOSITES AT 77 K 375

INFLUENCE OF PLY CRACKS ON FRACTURE STRENGTH OF GRAPHITE/EPOXY

LAMINATES AT 76 K 385

ORTHOTROPIC ELASTIC CONSTANTS OF A GLASS- FIBER-CLOTH EPOXYMATRIX COMPOSITE 415

CALCULATED ELASTIC CONSTANTS OF COMPOSITES CONTAINING

ANISOTROPIC FIBERS 429

TECHNOLOGY TRANSFER

TECHNOLOGY TRANSFER PROGRAM 457

INTERNATIONAL COOPERATION IN RESEARCH TO DEVELOP RADIATIONRESISTANT LAMINATES 458

ACKNOWLEDGMENT 464

iv
SUMMARY

This report contains results of a research program to produce material property data that will

facilitate design and development of cryogenic structures for the superconducting magnets of

magnetic fusion energy power plants and prototypes. The program was conceived and developed

jointly by the staffs of the National Bureau of Standards and the Office of Fusion Energy of the

Department of Energy; it is managed by NBS and sponsored by DoE. Research is conducted at NBS and

at various other laboratories through subcontracts with NBS.

The reports presented here summarize the sixth year of work on the 1 ow-temperature materials

research program. Highlights of the results are presented first. Research results are given for

the four main program areas: structural alloys, weldments and castings, nonmetal 1 ics , and

technology transfer. Objectives, approaches, and achievements are summarized in an introduction to each program area.

The major portion of the program has been the evaluation of the low- temperature mechanical and

physical properties of stainless steel base metals, welds, and castings, with particular emphasis

on the nitrogen-strengthened stainless steels. Developmental steels with manganese additions are also under investigation. Codes and standards development received increased emphasis this year,

especially as related to low- temperature property measurements. Work has been done on the

standardization of industrial laminates for low- temperature applications and on the measurement of

nonmetallic composite properties at cryogenic temperatures . Modeling of the elastic and fracture

behavior of woven fiber- reinforced composites has been initiated this year.

Note : Certain commercial equipment, instruments, or materials are identified in this paper to specify the experimental procedure adequately. 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. Papers by

non-NBS authors have not been reviewed or edited by the NBS editors. Therefore, the National

Bureau of Standards accepts no responsibility for the comments or recommendations contained

therein.

v
ORGANIZATIONAL CONTACTS

The people responsible for the various major aspects of the program are listed below so that specific technical questions may be directed to them.

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

Program Monitor D. Beard (301) 353-4958

(FTS) 233-4958

National Bureau of Standards, Boulder, Colorado 80303

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

(FTS) 320-3870

Structural Alloys R. P. Reed

R. L. Tobler

J. M. Arvidson

R. P. Walsh

(303) 497-3870

(303) 497-3421

(303) 497-3424

(303) 497-3537

Elastic Properties H. M. Ledbetter (303) 497-3443

Welding and Casting H. I. McHenry

Y.-W. Cheng

P. Purtscher

(303) 497-3268

(303) 497-5545

(303) 497-3699

Nonmetal 1 ics M. B. Kasen

H. M. Ledbetter

R. D. Kriz

(303) 497-3558

(303) 497-3443

(303) 497-3547

Technology Transfer N. J. Simon (303) 497-3687

Colorado School of Mines, Metallurgy Department, Golden, Colorado

E. L. Brown (303) 273-3953

Rockwell International Corporation, Energy Systems Group, Rocky Flats Plant,

Golden, Colorado

T. A. Whipple (303) 497-4365

University of Colorado, Mechanical Engineering Department, Boulder,

Colorado 80309

S. K. Datta (303) 492-7750

vi
PROGRAM DESCRIPTION

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

construction, and safe operation of 1 ow- temperature magnetic fusion

energy (MFE) systems, especially superconducting magnets, through effective materials research and materials technology transfer. The specific

steps taken to achieve this objective are: (1) evaluation of low-temperature materials research needs specific to MFE devices; (2) development and monitoring of a research program to acquire the necessary data;

and (3) rapid dissemination of the data to potential users through personal contacts, publications, and workshops.

Efforts directed at the first specific objective began with the

publication of the "Survey of Low Temperature Materials for Magnetic

Fusion Energy" in March 1977. A recent publication updating part of

this survey, "Structural Alloys for Superconducting Magnets in Fusion

Energy System," was included in Volume IV (1981) of this series. In

this volume, reviews of the properties of austenitic stainless steels

and of austenitic-steel elastic constants contribute to this objective.

Through continuous interactions with all low- temperature design, con- struction, and measurement programs, such as the Large Coil Project, we are aware of new problems as they arise. This year's contribution to

the second objective is described in Table 1 in the form of an outline

of the research projects. The results appear later in this report.

This year, more emphasis has been placed upon research oriented toward

establishment of test standards. The third objective is satisfied, in

part, by these annual reports and by the series of NBS-DoE Workshops on Materials at Low Temperatures, which are held biennially in Vail,

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

specific materials have been distributed to interested members of the

low- temperature materials community.
Table 1. Outline of the NBS/DoE Program on Material Studies for

Magnetic Fusion Energy Applications at Low Temperatures.

Program Area Organization Program Description

A. Structural Alloys

1. Development of strong, NBS

tough alloys

2. Development of strong, NBS

fatigue-resistant alloys

3. Development of low- NBS

temperature codes and

standards

Assessment of low -temperature properties of selected

austenitic alloys containing

varying concentrations of C,

N, and Mn with a Fe-Cr base.

Assessment of Japanese aus- tenitic alloys with Fe-Mn

base.

Assessment of fatigue and

fatigue crack-growth rates

of candidate structural

alloys, including 300 series

steels, Fe-Mn base alloys,

and age-hardened alloys.

Research to establish lowtemperature mechanical test

standards, now emphasizing

tensile, fracture toughness,

and fatigue.

B. Weldments and

1. Austenitic-steel NBS

wel ds

Castings

Evaluation of the strength

and toughness properties of

the 25Mn-5Cr and 32Mn-7Cr

alloys at 4 K. Continuation

of studies to understand the

mechanisms by which residual

delta ferrite affects the

deformation and fracture of

stainless-steel welds.

vi i i
Program Area Organization Program Description

2. Stainless steel

castings

3. Testing program

NBS, Colorado Investigation of the metalSchool of Mines lurgical factors that affect

the mechanical properties of

stainless-steel castings at

cryogenic temperatures, including metal 1 ographic

observations of duplex aus- tenite/del ta-ferrite microstructures in deformed and

fractured stainless steels.

Study of the effects of

ferrite and nitrogen con- tent on the strength and

toughness at 4 K.

NBS Continuation of mechanical

property testing at 4 K of

welds and plate material in

support of superconducting

magnet design and construction.

C. Nonmetal lies for

1. Standardization NBS

Magnet Structures

Continuing cooperation with

industry to establish and

refine cryogenic- grade specifications for insulating

laminates and coding systems

for laminates.

2. Characterization NBS Measurement of thermal

conductivity of composite

components at cryogenic

temperatures to aid in

establishing models for

predicting the thermal

conductivity of a variety of

laminates. Investigation of

failure modes of laminates

under compression at

cryogenic temperatures.

IX
Program Area Organization Program Description

3. Property determination NBS

model i ng

Development of dynamic methods to determine elastic

properties of laminated com- posites. For a woven glassfiber/epoxy composite,

experimental determination

of the complete ninecomponent elastic-constant

tensor. Study of the localstress distribution in a glass-fiber/epoxy composite

"unit cell" using finiteelement methods.

D. Technology Transfer

1. NBS/DoE Workshop NBS Presentation of research

results to the fusion com- munity, discussion of new

problems, and interaction

between interested parties

by means of a biennial

workshop.

2. Handbook NBS Distribution (1982) of

preliminary handbook data

pages on mechanical and

physical properties of AISI

316 stainless steels.

Distribution this year of

21Cr-6Ni-9Mn stainless steel

and AISI 304N and 304LN

stainless steels data pages.

Inclusion of aluminum alloy

data in future handbook

pages.

x
HIGHLIGHTS OF RESULTS

1

HIGHLIGHTS OF RESULTS

STRUCTURAL ALLOYS:

Tensile, Fracture, and Fatigue Properties

The studies leading toward development of strong, tough structural alloys for superconducting-magnet use continued, resulting in the following major accomplishments:

1. The temperature dependence of the flow strength, o^, of austenitic steels was shown to be

best characterized by a linear dependence of the log a

on temperature between 4 and

300 K.

2. Extensive reviews of mechanical properties, martensitic transformations , and elastic

properties of austenitic stainless steels were presented.

3. Little effect of magnetic fields on stainless steel mechanical behavior was identified.

Studies are continuing.

An extensive program to measure fatigue crack growth rates at low stress intensity factors was initiated. Highlights are:

1. An improved measurement capability was developed, including very sensitive compliance

correction terms and computer-interactive operation and analyses.

2. Elastic-modulus corrected measurements of aluminum, titanium, and stainless steels were

obtained, adding to calibration of measurements of crack growth.

Codes and Standards

Discussions were initiated with the Metals Properties Council to set up a cooperative program

leading to the establishment of low- temperature codes and standards. Working groups have been

organized into two areas: metals and nonmetals.

Elastic Constants of Fe-5Cr-26Mn, 76-400 K

By measuring longitudinal-mode and transverse-mode sound velocities at frequencies near

10 MHz, we determined the complete engineering elastic constants--bul k modulus, shear modulus.

Young's modulus, Poisson's ratio--for an Fe-5Cr-26Mn austenitic steel between 76 and 400 K. As the

result of a magnetic transition, all elastic constants behave anomalously below about 360 K. The

bulk modulus begins to soften during cooling at some higher temperature. All elastic constants

except Poisson's ratio resume an apparently normal temperature dependence below the 360- K magnetic

transition. After increasing abruptly at the magnetic transition, Poisson's ratio increases with

decreasing temperature.

3
For physical properties, these results imply important consequences. For example, properties

such as specific heat and thermal expansivity should also exhibit anomalies near the magnetic phase

transition. The overall 300-to-4-K thermal contraction should be much smaller than in conventional

304 stainless steel. Depending on their mechanism, plastic mechanical properties may also behave

anomal ously.

Austenitic-Steel Elastic Constants at Low Temperatures

Ledbetter reviewed his contributions to this topic: altogether 31 manuscripts, many un- published. Main sections of the review are: variability, alloying, temperature, monocrystal and

polycrystal relationships, texture, instability, and binary alloys (Fe-Ni).

WELDMENTS AND CASTINGS:

Austenitic-Steel Weldments

The tensile properties, fracture toughness, and fatigue crack-growth rates of two high-manganese austenitic steels were measured at 4 K. Base plate and weld metal properties were measured

in a 25Mn-5Cr alloy, and the base plate properties were measured in a 32Cr-7Mn alloy. The results

for both alloys fall near the strength-vs . -toughness trend line developed for austenitic stainless

steels.

Austenitic-Sta ini ess-Steel Castings

The tensile properties and fracture toughness of a series of five CM8M castings with nitrogen

contents varying from 0.02 to 0.20 percent were evaluated at 4 K. The nitrogen and chromium

contents of the alloys were varied within the limit of 9 to 11 percent nitrogen and 19 to 22

percent chromium, such that the delta ferrite content remained nearly constant at 9±1 percent. The

strength, ductility and toughness did not vary appreciably for nitrogen levels between 0.02 and 0.1

percent. The alloys with higher nitrogen contents had significantly higher yield strength, but

lower fracture toughness.

The metallography of CF8M castings with ferrite contents ranging from 1 to 28 percent was studied to determine the influence of delta ferrite on deformation and fracture. Evidence was found for two failure mechanisms: (1) brittle fracture of the delta ferrite mixed with ductile

fracture of the austenite, and (2) ductile failure in the austenite near the delta-ferrite interface.

Stainless-Steel-Weld Elastic Constants

For shielded-metal -arc E316L welds, we determined ul trasonical ly the complete set of engineering elastic constants: Young's modulus, shear modulus, bulk modulus, Poisson's ratio. Reflecting

texture, the elastic constants depend strongly on direction. Agreeing with many previous studies,

we found a strong <100>-axis texture perpendicular to the welded plate, in the principal heat-flow
direction. Disagreeing with these studies, we found that another strong texture element must be

involved to explain the measurements. We considered carefully the relationship between the macro- scopic weld constants and those of the microscopic constituent single crystals. Surprisingly, we found that delta ferrite up to 10.1 volume percent affects the weld elastic constants only negligibly.

NONMETALLICS

Standardization

NBS personnel continue to work with those in the United States laminating industry to provide

magnet builders with commercial sources of reliable, characterized nonmetal lie structural and

insulating materials. Of the two products in commercial production, G-10CR and G-11CR, primary

emphasis has been placed on the lower cost G-10CR, refining its specification and characterizing

its performance. Preliminary test data had suggested that the higher cost G- 1 ICR product

provided little advantage over G-10CR. However, more recent tests at ORNL have indicated that the

aromatic amine cure used in G- 1 ICR does provide better radiation resistance than the acid anhydride

cure used in G-10CR. Also, there is some evidence that the G- 1 ICR product may have superior

dimensional stability. Therefore, steps are being taken to bring the G- 1 ICR specification to the same level of development as the G-10CR specification.

Characterization

G-10CR products from five manufacturers were tested at 295 and 76 K, and the results were statistically analyzed to determine how much variability could be expected among manufacturers

.

Data on tensile strength, Young's modulus, compressive strength, Poisson's ratio, and interlaminar

shear strength were obtained. The analysis indicated that differences in performance were not statistically significant, suggesting that the specifications are adequate to insure uniformity of

product behavior.

A continuing study of the reasons for low interlaminar shear strength of a polyimide-matrix

variant of G-10CR has shown that one of the reasons is poor interfacial bonding between the E-glass

fibers and the matrix. Results of a fractographic study comparing the failure modes of the

standard epoxy product with that of the polyimide variant were presented at the 1982 ICEC9/ICMC

conference in Kobe, Japan.

Property Determination and Modeling

(a) Predicted elastic constants of composites containing anisotropic fibers . By a wave- scattering method, we derived dispersion relationships for waves propagating perpendicular to con- tinuous fibers that are oriented unidirectional ly. In the long-wavelength limit we obtained re- lationships that predict the composite's effective static elastic constants. We compared these

relationships with others derived from energy methods to obtain upper and lower bounds of the

effective static moduli. We plotted the predicted composite constants for graphite-epoxy over the

5
full range of fiber volume fractions. We consider the fibers to be anisotropic, but transversely

isotropic. Under special conditions, the energy-method upper and lower bounds compare identically

with the results of this study. The static properties are, of course, special cases of the more

general dispersion relationships. Graphs are given for nine elastic constants: axial and trans- verse Young's and shear moduli, bulk and plane-strain-bulk moduli, and three Poisson's ratios.

(b) Orthotropic elastic constants of a glass-fiber-reinforced epoxy-matrix composite . For a G-lOCR-type composite, we determined ul trasonical ly the complete nine-component orthotropicsymmetry elastic-constant tensor. Using conventional crystal-physics notation developed for single

crystals, we described the elastic constants both as Voigt C^.'s and S^.'s and as engineering

elastic constants: Young's, shear, bulk moduli, and Poisson's ratio. For the first time, anisotropic elastic strain and stress calculations can be carried out confidently for these materials.

(c) Prediction of local stress in a woven-composi te "unit cell." Woven-fabric composites are commonly used in superconducting magnets and in containment of cryogenic liquids. The mechanical

response of a plain-weave laminated composite at cryogenic temperatures was studied by predicting

the 1 oad-deformation response of a fundamental "unit cell." A finite-element model was used to predict three-dimensional stresses and elastic stiffness within the unit cell. Elastic Young's

modulus predicted by the finite-element model agreed exactly with that predicted by thin-laminateplate theory. From parametric studies it was found that the Young's modulus in the thin-laminate

plane is reduced 50 percent when the warp fiber bundle orientation is increased from 0° to 45°.

Parametric studies also demonstrated that thermal loads induced at 77 K decrease delamination

stresses at the fill-warp interface and increase delamination stresses in the matrix region adjacent to the warp fiber bundle. Decreasing warp fiber bundle angle decreases these stresses.

Future studies will include stress redistributions caused by fracture of fill fiber bundles.

TECHNOLOGY TRANSFER:

Handbook supporting documentation pages have been prepared for AISI 316 stainless steels and

distributed in preliminary draft format. Properties covered include: tensile strength, yield

strength, elongation, reduction of area, stress-strain curves, elastic constants, impact energy,

fracture toughness, stress- and strain-controlled fatigue life, fatigue crack growth rate, and

electromagnetic and thermal properties. These pages were also submitted to the advisory committee

of the "Materials Handbook for Fusion Energy Systems" and are currently in the review process.

Handbook and supporting documentation pages on 21Cr-6Ni-9Mn stainless steels have also been prepared, covering the same properties. The 21Cr-6Ni-9Mn pages have recently been submitted to the

advisory committee. Preliminary work on handbook pages covering AISI 304N and 304LN is in progress.

6
STRUCTURAL ALLOYS

7

STRUCTURAL ALLOYS PROGRAM

LEADER: R. P. Reed, NBS

STAFF: R. L. Tobler, J. M. Arvidson, R. P. Walsh, H. M. Ledbetter, NBS

OBJECTIVES:

(1) Development of strong, tough structural alloys for use at 4 K in superconducting magnets.

(2) Development of strong, fatigue-resistant structural alloys for

use at 4 K in pulsed superconducting magnets.

(3) Development of effective codes and standards related to lowtemperature property measurements and structural design.

RESEARCH PAST YEAR (1982):

(1) Review of martensitic transformations in austenitic steels.

(2) Characterization of Japanese Fe-Mn base austenitic steels.

(3) Preliminary development and review of low- temperature fatigue

measurements.

(4) Reports on effects of magnetic fields on austenitic-steel

properties.

(5) Characterization of low- temperature flow strength of austenitic

steel s. (6) Suggestions of modification of ASTM tensile specifications for

testing at 4 K. (7) Continued studies on low- temperature elastic properties of

austenitic steels.

RESEARCH THIS YEAR (1983):

(1) Continued research on effects of magnetic fields on 4-K

deformation of austenitic steels.

(2) Construction of 500 000-N ( 110,000-1 bf) capacity tensilefatigue testing facility for the temperature range 4 to 300 K. (3) Characterization of selected Japanese, United States, and USSR

Fe-Mn and Fe-Mn-Cr base austenitic steels.

(4) Correlation of metallurgical effects with fatigue resistance of

AISI 304, 316, and 310.

(5) Study of tensile strength and fracture toughness of manganesemodified 304LN austenitic steels.

(6) Development and verification of fatigue test methodology.

(7) Continued studies on low- temperature elastic properties of

austenitic steels.

(8) Thermal-expansion studies on austenitic steels.

9
SUMMARY OF MECHANICAL PROPERTY MEASUREMENTS OF STRUCTURAL ALLOYS

Alloy

(Designation)

Country

of

Supplier

Tensi le 295,76,4KElastic

295-4K

Fracture

Toughness

295,76,4K

Fatigue Crack

Growth Rate 295,76,4K

Fatigue

S-N

4K

Report Volume*

(page)

AUSTENITIC STEELS

Fe-19Cr-9Ni

(304)

U.S. MI,

III, IV,

1 ,11,

III, IV 1,11,111,IV

II, III, IV I 1(15,71,213),

11(79,149,175),

111(15,91,105,

1 17) , IV(37 , 101

,

203,215,227) , V ( 7 1

)

Fe-19Cr-9Ni

( 304L

)

U.S. 11,111,

IV Ml,IV III, IV III, IV 1(213), 11(79, 123,

175), 111(15), IV(37,

101,131,302,215)

Fe-19Cr-9Ni -N (304N)U.S. 1 , 1 1 1

,

IV

IV I, III, IV I, III, IV 1(93), 111(15), IV (37,101,203,215)

Fe-19Cr-9Ni-N

(304LN)

U.S. II, III,

IV,

IV 11,111,

IV

II, III, IV II 11(35, 79), 111(15),

I V( 37 ,101,203) , V( 29)

Fe-19Cr-9Ni

1 to 8Mn

( 304LN

)

U.S. IV,

V IV,I V( 77) , V( 15 ,59 , 189)

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

Fe-17Cr-13Ni-2Mo

(316)

U.S. Mi,

III, IV 1,11 II I 1(15, 71), 11(79),

111(49,105,117),

IV( 147) ,V( 185)

Fe- 17Cr-13Ni -2Mo

(316LN)

U.S. II II II 11(79)

Fe-25Cr-21Ni

(310)

U.S. V III V 111(105, 117), V( 71

,

145)

Fe-18Cr-3Ni-13Mn U.S. I III I I 1(93), 111(91)

Fe-21Cr-6Ni-9Mn

(21/6/9)

U.S. I III I I 1(15, 71), 111(91)

Fe-21Cr-12Ni-5Mn U.S. 1,11 III Mi MI 1(9?), 11(79),

111(91)

Fe-13Cr-19Mn USSR III 111(79)

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

Fe-19Ni-9Co

(200-300 grades)

U.S. IV I V ( 237

)

Many al 1 oys Review

IV,

Review

IV,

Review

IV,

Review Review

IV, V IV IV( 17,257) ,V( 171)

ALUMINUM ALLOYS

Al -4Mg

(5083-0)

U.S. IV II 11(35), I V( 185)

Al-6Cu-0.2Mn-0.1Fe

( 2219-T87

)

U.S. II II II II 11(19)

SUPERCONDUCTORS

Nb-45TU.S. III 111(133)

*Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures : NBS1R 78-884 (I);

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

10
MARTENSITIC TRANSFORMATIONS IN Fe-Cr-Ni STAINLESS STEELS+*

R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Very 1 ow- temperature martensitic transformations are of great concern for

cryogenic applications and research. The principal transformation characteristics are reviewed and then elaborated. The materials classes or alloy

systems that exhibit martensitic transformations at very low temperatures are

-briefly discussed, and references to reviews and compilations are given. The

austenite stability of Fe-Cr-Ni alloys with regard to cooling, elastic stress,

and deformation is discussed in detail, and the empirical stability expressions that have been developed are summarized and reviewed. Structural

relationships between the face-centered cubic austenite (y) and body-centered

cubic (a') and hexagonal-close-packed (e) martensite products are examined.

Typical stress-strain curves of stable and metastable austenitic alloys are illustrated and analyzed as is the temperature dependence of the flow strength

below room temperature. Finally, the influence of martensite transformation

upon alloy performance in cryogenic applications is surveyed.

+ Work supported by Office of Fusion Energy, Department of Energy

* To be published in Austenitic Steels at Low Temperatures , Eds. R. P. Reed

and T. Horiuchi, Plenum Press, New York (1983).

11

INTRODUCTION

The AISI 300 series stainless steels, especially alloys 304,

310, and 316, are used extensively for cryogenic structures. These

alloys have a high elastic modulus, high toughness and ductility, low

thermal and electrical conductivity, and good weldability. Because

of these characteristics, they are used more than any other alloy

class for structures for service at 20 K and below, but their use is

complicated by the metastability of the austenitic structure of most

alloys of this Fe-Cr-Ni alloy series. The metastability leads to

martensitic transformation during cooling, from applied stress or during plastic deformation. The martensitic transformation is a significant design consideration in applications requiring fracture

control planning, close dimensional tolerances, the absence of a ferromagnetic phase, and high toughness of weldment, heat-affected

zone, and base metal. Consequently, research has been conducted to

characterize the transformations and their effects on mechanical and

physical properties and in-service performance.

Martensitic transformation initially meant the transformation

from face-centered-cubic (f.c.c.) austenite to body-centered-cubic

(b.c.c.) or body-centered-tetragonal (b.c.t.) martensite in steels.

Gradually, however, metallurgists began to describe newly discovered

nonferrous transformations as martensitic in nature, and the term

martensitic evolved to describe a class of structural transformations. Cohen et al., 1 who have presented a thorough discussion of

the characterization of martensitic transformation, state "A

martensitic transformation is a lattice-distortive, virtually

diffusionless structural change having a dominant deviatoric component and associated shape change such that strain energy dominates

13
the kinetics and morphology during the transformation." Thus, a change of atomic structure that takes place'with no atomic diffusion

and exhibits a change of shape with an attendant undistorted plane is

considered a martensitic phase transformation. In such transformations, the associated strain energy is very significant in the determination of the final product shape and size and in the transformation time and temperature characteristics.

Nucleation restricts the martensitic transformation in all

ferrous alloys and in some nonferrous systems. After nucleation,

growth to a restricted product morphology usually proceeds very

quickly; the rate may approach the speed of sound. Except in certain

nonferrous alloys, additional transformation is achieved by the nu- cleation of new crystals, not by the growth of the existing crystals.

Martensitic transformations occur at very low temperatures in

such diverse materials as stainless steels, alkali metals, and solidified gases; therefore, atomic diffusion plays no essential role.

Atom transfer from parent to product structures proceeds by shearlike, coordinated motion of many atoms, resulting in a shape change.

This shape change is easily observed by using conventional microscopy

techniques

A consequence of the shear-like, coordinated atomic motion is

that there are easily defined geometrical relationships between

parent and product phases. An orientation relationship exists between all product-phase and parent-phase crystal structures. The

interface plane, separating the parent and product phase, is conventionally called the habit plane and is usually unique for a given

alloy system. Phenomenological crystallographic theories have been

reasonably successful in predicting and explaining the interrelationships of these geometrical parameters, but they are not completely

satisfactory. Another consequence of the atom-transfer mode of a martensitic transformation is that the identical atom concentration

of the parent phase is preserved in the .product phase; there is no composition gradient between the two phases. Finally, locally large

stresses, which may exceed the yield strength of the parent material,

may result from the shape deformation of the product phase. This is

particularly apparent in ferrous-alloy transformations, where extensive slip, faulting, or in special cases, deformation twinning may be

induced in the parent phase and where the generation of point defects

in the product phase may be about 0.5 at.%.

No review has previously been devoted to martensitic

transformations in Fe-Cr-Ni steels. However, here are a number of

excellent reviews on martensitic transformations. Especially, the

reader is referred to Reed and Breedis 2 for a thorough compilation

(702 references) of literature prior to 1966; to Bilby and

Christian, 3 Christian, 4 Wayman, 5 ’ 6 and Lieberman 7 for excellent

descriptions of martensite crystallography; and to Roitburd and

14
Kurdjumov, 8 Entwisel, 9 Christian, 10 and Kaufman and Cohen 11 for

excellent reviews of martensite kinetics. Low temperatures have been

emphasized in reviews by Barrett 12 and Reed and Breedis. 2 General

books on martensitic transformations include Martensitic Fundamentals

and Technology 13 and Christian. 14 There have been excellent international conference proceedings on general martensitic research:

Physical Properties of Martensite and Bainite , 1

5

Mechanism of Phase

Transformations in Crystalline Solids , 10 Shape Memory Effect in

Alloys , 1 ^ New Aspect of Martensitic Transformations , 10 and Phase

Transformations . The International Conference on Martensitic

Transformations (ICOMAT) held its third conference in 1979. These

ICOMAT conferences bring most of the principal research scientists

together every third year. Their proceedings (e.g.. Proceedings of

the International Conference on Martensitic Transformation,

ICOMAT-1979 , 2u report current research.

AUSTENITE STABILITY

Either on cooling, under applied elastic stresses, or during

plastic deformation, the austenite phase (y) of Fe-Cr-Ni alloys may

transform martensitically to b.c.c. (a 1 ) and h.c.p. (e) phases. The

alloy composition affects the chemical free-energy difference between

the two phases.

The difference in free energy (F) between the two phases y and

is

AFa'+y

( 1 ) Y a where F^ and F^ are the total free energies of tl?e respective

phases. The free energies of mixing, AFT and AF , of the two phases

may be expressed in terms of the total free energies as:

afy = Y o

: Fi, - EX. F. and (2) m T li

a'

AFm

= f“' -

Zx. F°

T li (3)

where X. are the atom fraction of a component, i, and F? is the free

energy of the pure component, i, in the standard state. The total

free energy of each phase is a sum of the partial molar free energies

of each component in an ideal solution of y or a':

fI = ex. fT

a ' -a

and F = EX. F.

T li (4)

The difference in the free energy of mixing for the two phases is:

AFl

- AF“' = Fl - EX F° - f“' + EX. F? and (5) m m T 11 T 11

Y a AF' - AF m mY a* a’-^Y = F

Y - F

t = AF T

( 6 )

15
( 7 )

Therefore, from eqs. 4 and 6,

AF

a

= EX.(fT - F

a ') = IX. (AFa

'^

Y

)

That is, the free energy difference between the two phases can be

expressed as the sum of the partial molar free energy differences of

each component.

oi a' -*-y Breedis and Kaufman*11 have derived the expression for AF for

the Fe-Cr-Ni ternary system:

“ (1 - X

Cr

-a +Y ,, . -a

iAF„ + X_ AF^ 1 + ^i^'Fe Cr Cr

X

N1 )(B - A) Fe-Cr

+ (8)

X

N1

(1 - X

Cr - V(B - A) Fe-Ni

+ X

Cr

X

Ni

(B ‘ A>Cr-M

x„.af“/y + X- (1 - X. Ni Ni Cr Cr

Here X and X are the atom fractions, and A and B are interaction

parameters. Tnese parameters are described as

exc exc A = X ( 1 - x) and B = X ( 1 - x) (9)

Ot

' ->y wheye F is tfye excess free energy of a binary system. AF , AF^ Y

, 6

and ^Fp

e Y represent the partial molar free-energy difference

between a' and y for the pure components chromium, nickel, and iron,

respectively.

Kaufman 22 derived free-energy relations for the Fe-Cr binary

system and combined these,with the previous data for Fe-Ni binary

alloys 23 to calculate AF

U Y for stainless steels. On the basis of

regular solution and for the temperature range of about 100 to 900 K,

AfJTY = AFY - AF

a Fe,Cr,Ni

= (1 - X

Cr - XNi )(5036 - 11 x 10" 3 T

2 + 6.5 x 10

_6

T

3

)

+ X

Cr

(1927 + 4.2T) + X

N±

(-15500 + 2.97 x 10

_3

T

2 + 1.64 x 10

_6

T

3

) +

Xr (1 - X_ - X ) (-1 1730 + 3. IT) + (10) Cr Cr Ni

XNi (1 ~ X

Cr " ^i^ 15080 + 2 *4T(l-JlnT)] J*mol

_1

For calculations of T of stainless steels, Kaufman22 added the term

ms

X„ (-2346 + 0.175T) J-mol"

1 (11) G+N

to account for carbon and nitrogen; the nominal manganes?_content of

1.5 wt.% was added tothe nickel content. Values of AF

a Y ranging

from about 2000 J»mol (475 cal*mol ) for T = 100 K to 1260 ms

16

i
J«mol * (300 cal«mol *) for T = 550 K were calculated for typical

stainless steel compositions using Eq. 8. (In this paper the temperature of transformation start is labeled T ; this nomenclature is a

IDS compromise between the metallurgical expression, M

g

, and the physics

label, T . ) Data calculations bj£, Breedis and Kaufman21 and Rao et

al. 24 predict lower, ^alues of and would, similarly, result in

lower values of AF

a Y for the Fe-Cr-Ni ternary system.

An equation analogous to Eq. 10 can be developed for the free

energy difference between the e and y phases in Fe-Cr-Ni alloys.

Values for AF^ Y and AF*lTY were estimated by Kaufman25 from enthalpy

and entropy differences due to structure. Interaction parameter es- timates were made by Breedis and Kaufman21 for Fe-Cr, Fe-Ni, and

Cr-Ni binaries. Using these estimates, the following expression is

derived for the free energy difference between e and y phases in

Fe-Cr-Ni alloys:

AFen = X AF

£ *Y

Fe,Cr,Ni A

Fe *Fe

+ X

Cr

(1095 + 0.628T) - XNi (1047 + 1.26T) +

2095 (XCrX

Cr

+ ^Ni V J*mol

-1

( 12 )

The partial molar free-energy difference (AF^ ) can be calculated from the graphical presentation by Breedis and Kaufman21 of

AF , and AF . This term becomes positive at temperatures lower

than 390 K andequals about 300 J*mol at 300 K. Using Eq. 12 and

the^AFp

e Y data plotted in Fig. 1, the temperature dependence of

AF^, Y 6

. can be calculated and is plotted in Fig. 1. This indicatlS Efta£ AFp

g

Y

m-[ becomes positive near 325 K. Actually, e transformation has’ been observed on cooling below room temperature

(e.g., Reed 26 ). Alloy additions such as carbon, nitrogen, and manganese move the AF

e+y

Fe,Cr,Ni curve to lower temperatures.

Expressions relating the stability of the austenite and the temperatures of transformation during cooling (T ) or during deformation (T ) have been developed empirically an9S are summarized in

Table 1. Eichelman and Hull, 27 working with austenitic alloys with

10 to 18% Cr, 6 to 12% Ni, 0.6 to 5% Mn, 0.3 to 2.6% Si, 0.004 to

0.129% C, and 0.01 to 0.06% N, established that all of the above

alloying elements stabilize the austenite and thus lower T

Monkman et al., 28 using a larger number of specimens with ranges of 5 to 13% Ni, 11 to 19% Cr, and 0.035 to 0.126% C+N, produced a similar

analysis and concluded that only to a first approximation was T

linearly dependent on composition. The dependence of T on carfon

and nitrogen concentration seemed to be influenced by chromium and

nickel concentration.

Hull 31 completed another study assessing the effects of nickel,

chromium, manganese, carbon, nitrogen, silicon, and cobalt (aluminum,

thallium, vanadium, and tungsten sometimes were added) on a'

17
Fig. Temperature dependence of free energy difference of

e and y structures of iron and Fe-18Cr-8Ni alloy.

formation during low-temperature cooling-and-def ormation experiments

and on ferrite retention after cooling to room temperature from the

melt. Hull's study reconfirmed that all the above elements suppress

low-temperature a' formation during either cooling or deformation but

that chromium, molybdenum, silicon, vanadium, tungsten, niobium, tantalum, and aluminum additions promote the formation or retention of

the high-temperature, b.c.c. delta-ferrite phase. To predict the

effect of elemental additions on T , Hull assumed that the nickel

contribution was the average of the Monkman et al. 28 and Eichelman

and Hull 27 results and then compared all the effects of alloying

additions with that of nickel. The empirical relationships have been

analyzed by King and Larbalestier

,

140 who point out some expected

deviations in the dependence of T on nickel and chromium when the

relationships are used for prediction of the composition range of

AISI 300 series stainless steels.

The formulation of Andrews 30 applies to lower chromium and

nickel concentrations, thus higher T , and his results are included

ms

18
in Table 1 to point out the disparity between the high- and lowtemperature empirical results. Hammond 29 was the only investigator

to include molybdenum with chromium and nickel.

In all these studies, specimens of various composition were

prepared and cooled, and their T values were measured. Regression

analyses were then used to obtain the empirical relation of the dependence of T on alloy concentration.

Similar experiments have been conducted to assess the influence

of cold work (either tensile or compressive) on T . Their results

are also contained in Table 1. Note that the experimental definition

of T varied in each study: Angel 32 used 50% a* at 30% tensile

elongation; Williams et al. 33 used 2.5% a' at 45% compressive elongation; and Hull 39 used the minimum detection level (probably about

1 % a') at 50% compressive elongation.

In reviewing the current status of the predictability of

Fe-Cr-Ni stability, Collings and King91 caution that the Breedis and

Kaufman21 free-energy expression for T near 0 K does not have the same chromium and nickel coefficients as the Eichelman and Hull 27 and

Monkman et al. 28 T expressions. The latter expressions are based

predominantly on chromium and nickel concentrations less than those

required for T near 0 K (T ^ 300 K) . This raises the question of

the usefulness of these T predictive expressions for the normal

18Cr-8Ni stainless steels?S

Collings and King91 also point out that in low- interstitialaustenitic stainless steels (those steels that are most easily

welded), there is a higher tendency for isothermal transformation (in

the 100 to 200 K range), and these predictive equations are not

useful in the estimation of the occurrence of time-dependent a'

transformation (or the separation between athermal and isothermal

transformation)

.

Schramm and Reed 37 used x-ray peak-shift measurements, coupled

with previous stacking-fault energy measurements, to estimate the dependence of the stacking-fault energy on chemical composition. This

dependence should correspond to the dependence of the y+e transformation on composition, because a stacking fault represents a local

planar area of e. Rhodes and Thompson 38 suggested that the stackingfault energy values of the Schramm and Reed analysis were too large,

considering additional weak-beam electron-microscopy data. Weak beam

electron microscopy measurements by Bampton et al. 92 confirm that the

Schramm and Reed least-square analysis produced stacking fault energies that were excessively large and indicated a data spread of about

±25% from measurements of individual nodes. Brofman and Ansell 39

added the dependence on carbon. The stacking-fault-energy composition dependences are also included in Table 1. The stacking-fault

19
Table 1. Temperature equivalents for calculation

Investigator (Year) Base Cr Ni

Temperature Equivalent

Mn Si C

f .c.c.+b.c.c. (T ), cooling3 ms Eichelman and Hull 27 1578 -41.7 -61.1 -33.3 1 ro 00 -1670

Monkman et al.‘ 8 1455 -36.7 -56.7 -146.0

Hammond'11105 -29 -39

Andrew 30 273 -12.1 -17.7 -30.4 -423

Hull 31 1755 -47 -59 -54 -37 -2390

f .c.c.**b.c.c. (T ,), deformation3 md

Angel 32 686 -14 -9.5 -8.1 -9.2 -46.2

Hull 31 1655 -23 -59 -41 -20 -777

Williams et al

3

686 -6 -25 -16 + 21 -222

Magnetic transition (paramagnetic+antiferromagnetic , N£el temperature)

Warnes and King 34 90 -1.25 -2.75 + 7.75 -14

Enthalpies of Formation*5

Lee et al

33 8168 208 181 1087 3186

f .c.c.+h.c.p. (stacking fault energy at room temperature) -

Dulieu and Nutting 36 0 5 14 34

Schramm and Reed 37 -530 7 62 32

Rhodes and Thompson 38 12 6 14 177 -47

Brofman and Ansell 39 167 9 21 260

a. Numbers in table represent K/wt.%; b. Numbers in table represent J*mol '‘wt.%; from enthalpy measurements, lath, and plate a'.

20
of stability parameters of austenitic steels.

N Mo Other

Comments, Composition Range

(wt.%)

-1670 21 alloys: 10-18Cr, 6-12N1, 0.6-5Mn, 0.3-26S1,

0.004-0. 12C, 0.01-0.06N

-1460 49 alloys: ll-19Cr, 5-13N1, 0. 035-0. 0176(C+N)

-36 16 alloys: 0-12Cr, 4-8Ni, 0.03C, 2-6Mo, 0-15Co,

1-2T1

-7.5 184 alloys from previous studies not in this table.

Notice different composition ranges. 0-4. 6Cr,

0-5. ONi, 0.04-4.9Mn, 0.1-1.9S1, 0.11-0.6C, 0-5. 4Mo

-3720 -56 -180(Ti)

-14 (Co)

, 59Ni = average of Eichelman and Hull 27 and Monkman

et al., 28 29 alloys: l2-24Cr, 0-22N1, 0-20Mn, 0-4S1,

0-0. 1C, 0-0. 15N, 0-6Mo, Co, 0-2T1

-46.2 -18.5 30% tension, 50% a'

-315 -24 -12(Co) 50% compression, 60 alloys: 12-24Cr, 0-22N1, 0-20Mn,

0-4Si , 0-0. 1C, 0-0. 15N, 0-6Mo, Co

-222 -11 45% compression, 2.5% a', 25 alloys: 12-25Cr,

9-20Ni, l-2Mn , 0.1-0.6S1, 0.04-0.25C, 0.01-0. IN, 0.6-2.8MO

-5.5 64 commercial and specially prepared alloys: 5-25Cr,

5-32N1, 0-16Mn, 0-24S1, 0.3Mo

113 7 d 48 alloys: 0-18Cr, 0-30N1, 0-3Mn, 0-0.05S1,

0.01-0.86C, 0 . 02-0 . 2Mo

1 36 (Ti) 32 (Cu)

-055 (Co)

, Electron microscopy

9

93 7 alloys, x-ray, literature data-regression analysis

Modified Schramm-Reed approach

c. Numbers in table represent pj/m2 *wt.% (ergs/cm2 *wt .%) ; d. Represents stored energy

21
energy is expected to decrease at lower temperatures, and the temperature dependence should approximately follow the form of the

AFf, ^ _ ... curve illustrated in Fig. 1. Fe.,Cr,Ni °

In all empirical relations, chromium is shown to increase the

stability of austenite relative to either e or a' (Table 1). Although it is not clear at first glance, Eq. 8 (intended for use below

900 K) predicts that chromium will act to suppress the y->a* transformation. At higher temperatures, chromium has the opposite effect and

stabilizes the y structure. 22 ^ However, Eq. 12 predicts that chromium

additions act to increase AFf,^^ thus promoting e transformation. This contrasts with tne ’ regression analyses results of the

dependence of the stacking-fault energy on composition (Table 1).

Additional study is needed to decide whether to amend the free-energy

expressions and whether chromium alloyed with manganese, silicon,

carbon, and nitrogen is a stabilizing influence on austenite with

respect to hexagonal-close-packed structures.

Warnes and King 34 measured the Ndel temperature (paramagneticto-antif erromagnetic transition temperature) for a series of Fe-Cr-Ni

alloys. Regression analysis of their data has resulted in an expression relating the Ndel temperature to composition (Table 1). The

magnetic term was considered to be a relatively minor contribution to

the free energy of the Fe-Cr-Ni alloy structures (e.g., Breedis and

Kaufman 21 ). It is of interest, however, that both chromium and

nickel act, in the approximately same ratio, to suppress the antiferromagnetic phase and to stabilize the austenite structure.

Calorimetric measurements of the heat evolved during martensite

formation were made by Lee et al. 35 for a series of Fe-Cr-Ni alloys.

For an Fe-18Cr-8Ni alloy transformed at temperatures ranging from 156

to 293 K, the transformation enthalpy change, ^AH, was measured and

the average value was found to be 1920 J*mol . No temperature

dependence of AH was observed. Results of the regression analysis of

all alloys for elemental contributions to the enthalpy change are listed in Table 1. There is a strong similarity between these re- sults and the effect of the contributions of the elements on Tms

STRUCTURE

Face-centered-cubic austenite (y) transforms on cooling to two

martensitic products: one b.c.c. (a’), the other h.c.p. (e). The

h.c.p. phase is associated with extended stacking faults and forms as thin sheets on (111) austenite planes. In most austenitic stainless

steels, the a* product forms as laths. Breedis 43 reported a lath a' morphology for compositions ranging from Fe-19Cr-llNi to Fe-lOCr16Ni; at lower chromium and higher nickel concentrations the a' mor- phology changed to a plate-like structure.

22
Fig. 2. Schematic representation of a' laths within a {111} sheet

in Fe-18Cr-8Ni alloy.

From Reed, 26 the a' lath-like structure is parallel to <110> , with (225) , (112) , or both habit planes. 44 The laths are re-

^ stricted within {ill} bands, and usually three sets of habit planes

form within the band.^ The schematic in Fig. 2 is a typical representation of a’ laths within a {111} band. For a given <110> lath direction, two possible habit planel are possible; for exampll, for

[110] (111) , both (225) and (225) are possible habit planes. The

habit^plane Observed is tKe one that^makes the largest angle to the

confining (111) band. Many adjacent laths appear to be twin

related. 35 Optical photomicrographs of bands of a' martensite, taken

in sequence after polishing and etching, are shown in Fig. 3. Reed 26

notes that during thermal cycling or cooling to a lower temperature,

new bands of lath martensite continued to form, and morphological

variations were observed. Some lath directions deviated from <110>,

and the habit plane deviated from {225}. While most laths were

confined by {111} bands, some a' was restricted by {211} bands.

Therefore, continued transformation adds considerably to the complexity of the a' structure.

Transformation during plastic deformation initially results in

a' at intersections of active {111} deformation bands. The a*

assumes the shape of laths along <110> common to the two active {111}

systems. Figure 4 illustrates a' formation at slip-plane intersections. Evidence of dislocation pileups at a* laths is observable.

Also, more than one lath may form at each transformation site, and in

contrast to the a' that forms on cooling (Figs. 2 and 3), all laths

have the same habit-plane variant. Finally, at several sites, the

lath transformation appears irregular or incomplete; there is a lack

23
I

•* 4

* ,

St % i

*

. < % ^

>

# ft <'

.'V

‘ 1

/

*

•**IIf

' f?**

v »*

l* ^ % x

4 &

* «

(a) a) ( c \ (<S) ( S /

Fig. 3. Optical photomicrographs of sheets of a' martensite in

Fe-18Cr-8Ni taken after sequential electropolishing and

etching steps (a to e).

of a well-defined austenite-martensite interface, and the interaction

of transformed material with the dislocations in the immediate vicinity is very complex. Growth of the a’ nucleus produces a more distinct and crystallographically regular lath-like product-phase mor- phology .

There is a considerable amount of accommodation deformation in

the austenite after transformation during cooling; it is particularly

large within the {111} band containing the a* laths. This leads one

to suspect that lattice deformation is minimized normal to the {111}

band that contains the three < 1 1 0 > directions and that the latticeinvariant deformation in these <110> directions may be very low. The

{111} bands contain either e martensite or a large amount of

stacking faults. 1+3 ,t+t+

The orientation relationships between the y, e, and a* phases

are:

24
(111) 1

1 (0001) I I (101) , Y

1 e1 a

[Uo]

Y

1

1 [1210]

e

|

1 [iii]

0

.

These relationships are apparently retained regardless of the manner

of e or a' formation.

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

consists predominantly of dislocations. Wirth and Bickerstaf f

e

4

5

have reported the presence of some twins in a' following increased

chromium additions (with 8 wt.% Ni) or increased cooling rates. As

chromium is replaced by nickel, Breedis 43 reports that cellular, irregular distributions of dislocations typical of a' laths change to

planar, regular arrays typical of a' plates. Also, as nickel re- places chromium, the amount of e transformation decreases, the sharpness of the h.c.p. reflections decreases, and the f.c.c. twin reflections become more diffuse. 43

Many metallurgists have been concerned with the role of the e transformation: does e act as a precursor transformation (y^e^a 1 ) or

is the e an accommodation effect (y+a', y+e)? Both effects have been

Fig. 4. Fe-18Cr-8Ni austenite strained 0.01 in tension with two_

active slip systems on {111}. Formation of a' along <110>

intersections of slip planes.

25
observed. The stainless steels with low stacking-fault energy tend

to form e, and the a' tends to form from e. Higher stacking-faultenergy alloys require transformation stresses for observable e formation .

EFFECTS OF TEMPERATURE

The amount of a* transformation is dependent on temperature, and

the kinetics generally follow the form shown in Fig. 5. After

cooling below T , holding at constant temperature produces a C curve when temperature is plotted against holding time for constant a' product. Suppression of a' transformation usually occurs below about

100 K. 46 In most of the Fe-18Cr-8Ni (AISI 304L) alloys that have

been examined and found to exhibit a' transformation, cooling to 76 K

produces a maximum amount of a’ (usually 5 to 10%). Cooling rapidly

to 4 K in some cases suppresses a’ transformation entirely.

Reversion of e, formed either from deformation26 or cooling, 47

to y, begins when the alloy is heated to about 410 K (see Fig. 6).

Reversion of a', formed either by cold working (^70% a ) , or cooling

(^>7% a'), begins at a higher temperature, about 770 K as illlustrated

in Fig. 26. 26 Partial decomposition of the a' on holding at lower

temperatures has been observed.

EFFECTS OF AUSTENITE DEFORMATION

Deformation of austenite at temperatures several hundred degrees

above room temperature tends to produce a cellular dislocation substructure. At lower temperatures (about one hundred degrees above

room temperature) more planar defects, stacking faults, and twins are produced. At still lower temperatures, e martensite, a* martensite,

or both form. The exact temperatures depend primarily on the

stacking-fault energy. The lower the stacking-fault energy, the

higher the deformation temperatures for formation of planar dislocations, stacking faults, and e martensite. Since the stacking-fault

energy is very dependent on temperature, 42 ’ 47 deformation at lower

temperatures results in an increase of planar defects.

The formation of planar defects tends to enhance subsequent martensite formation at lower temperatures. In contrast, a cellular

dislocation structure suppresses subsequent transformation on cooling

below T^^ or during deformation below T^. Presumably cellular

structures only influence the martensitic transformation by raising

the strength of the austenitic matrix. Planar defects, with reduced

cross-slip, produce high local strain fields, which may serve as nu- cleation sites. Evidence of austenitic prestrain effects on transformation at lower temperatures confirms the very delicate balance

between austenite strengthening and planar-dislocation-induced stress

fields

26
MARTENSITE,

percent

Fig. 5. Typical martensite-transformation-kinetics

for Fe-18Cr-8Ni alloy on cooling below room temperature. The series of curves show the

effect of progressively decreasing cooling

rates, from (a) to (e)

.

Small amounts of tensile plastic deformation (0 to 10%) at room

temperature stimulate transformation to a' on cooling to temperatures

below T < 26»46>48 Added deformation steadily suppresses transformation. Sreedis 48 has reported that the a' transformation is enhanced

in alloys that have a low stacking-fault energy; planar dislocation

arrays, as opposed to cellular dislocation substructures, result in

more transformation during cooling. Strife et al. 48 have found that

prestrain above T reduces T , may enhance subsequent transformations, and at larger strains, usually suppresses transformation.

There is ample evidence that deformation-induced defects play a major role in the nucleation of both e and a*. Brooks et al. 50

Fig. 6. Reversion of a' and e martensite on warming, following

formation at 76 K in an Fe-18Cr-8Ni alloy (AISI 304L)

.

27
concluded that the nucleus of the e martensite is a single stacking

fault and that a' nucleates from faulted regions in dislocation

pileups. Others (Suzuki et al., 51 Mangonon and Thomas, 52 Lagneborg,

Murr et al. 53

) have observed that the formation of strain-induced a' initially occurs at intersections of two active {111} slip systems

(see Fig. 4)

EFFECTS ON AUSTENITE DEFORMATION

Below the equilibrium temperature, T , martensite tends to form

during (a) cooling, (b) application of elastic stress, and

(c) plastic deformation of the austenite. Transformation during

cooling and during application of elastic stress (stress-assisted

transformation) results in similar martensite morphology. But, as described earlier, strain-induced martensite forms at slip-band intersections. During plastic deformation of metastable Fe-Cr-Ni austenite at low temperatures, two strain-induced martensitic phases

form, h.c.p. e and b.c.c. a'.

The AISI 300 series stainless steels used in cryogenic applications range from metastable to stable austenites. Alloys, such as

304, containing 18% Cr and 8% Ni, are metastable, whereas 310, con- taining 26% Cr and 20% Ni is stable with respect to martensitic

transformation. The stress-strain behavior and temperature dependence of the flow strength of these alloys differ and depend on the

austenite stability. In the more unstable alloys, both e and a' martensite are formed; in the slightly metastable alloys (e.g., 316),

only a* is formed; and in the stable alloys, neither a' nor e is

formed during deformation to fracture at any temperature.

Figures 7 and 8 assist in explaining the role of martensite

transformations in austenite deformation. In a metastable alloy at

low temperatures (4 to 200 K) , three distinct stages are present. A

stable alloy only exhibits dislocation work hardening and has "wellbehaved" stress-strain characteristics.

Stage I (Fig. 7) of metastable austenite strain-strain behavior

represents the microstrain and early macrostrain behavior that includes the 0.2% offset yield strength. The formation of a’ is not

thought to occur in this range; perhaps stacking-fault clusters, or

the e phase, or both contribute to low-temperature deformation in

this range. The stacking-fault energy is reduced at low temperatures, and apparently becomes low enough to promote e martensite.

With x-ray analysis to detect e formation during early deformation

and creep at low temperatures, Mirzagev et al. 54 related the log of

volume fraction of e over the volume fraction of y + a' to the log of

the plastic deformation for an Fe-18Cr-13Ni-0 . 02C steel; other x-ray

analysis measurements on similar steels have not confirmed this

e-strain correspondence. 54 ’ 55

28
~0.03 ~0.07 STRAIN

Fig. 7. Typical stress-vs. -strain curves of stable and metastable

austenitic Fe-Cr-Ni alloys at low temperatures and

associated relative amounts of strain-induced e and a* that

form in the metastable alloy.

The temperature dependence of metastable and stable Fe-Cr-Ni

alloys in Stage I is- illustrated in Fig. 8. There are three regions

of the temperature dependence of the flow strength for metastable

alloys: (a) the high-temperature range (>Tj) that exhibits normal

temperature dependence with no detectable martensitic products;

(b) the middle region (T < T < T^) in which the flow strength de- creases with decreasing temperature; and (c) the low-temperature

range (<T ) where flow strength again increases at lower temperatures. Tne temperature dependence of the flow strength of stable

alloys is normal, increasing monotonically with decreasing temperature .

In Stage I, a* has not been detected. Magnetic torsion-balance

measurements 56 do not sense a permeability change until Stage II is

reached. Suzuki et al. 51 have also used magnetic measurements to

characterize the stress level at which a' begins to form as a function of temperature; their results are also included in Fig. 8. The

stress level at which a* forms is considerably above the yield

strength. At first, the study of Olson and Azrin 57 appears to

conflict with these premises. They worked with a less stable

Fe-9Cr-8Ni alloy (a TRIP steel) and obtained an excellent correlation

between the temperature dependence of the yield strength and the flow

strength at which a' was first detected by using ac permeability measurements. It is likely, however, that martensite formation in their

alloy was stress assisted and that e did not form.

29
(9

Fig. 8. Temperature dependence of tensile flow strength at 0.2%

offset, 0

^ ^

and estimated strength at which a' begins to

form, aa ,,’of stable and metastable Fe-Cr-Ni steels. 51

Stress-assisted martensite formation or the formation of a* from

preexisting nucleation sites is best illustrated by the study of

Breedis and Robertson. 58 They measured the temperature dependence of

a metastable single-crystal Fe-16Cr-12Ni and of a stable singlecrystal alloy. The critical resolved shear stress of the metastable

alloy fell to zero at the T of a' (^240 K) . They detected the presence of both e and a' in tin! deformed crystals at temperatures above

T . The critical resolved shear stress of the stable alloy inms creased continuously with decreasing temperature. The T of commercially available alloys is normally below 0 K, and there¥ore, the

flow strengths of these alloys are not expected to approach zero.

Stage II, consisting of an "easy glide" range is best correlated

with the formation of a' laths at cross-slip intersections. Suzuki

et al. 51 have proposed that such a’ laths, with the long <110> direction representing the intersection of two active slip systems, act as windows to assist cross slip. Stage II is most prevalent in the temperature range 76 to 200 K. In this temperature region, the range of

strain at which the stress remains relatively constant is large,

usually about 0.03 to 0.05.

In Stage III, the rate of work hardening increases to a constant, which is maintained over a relatively large amount of plastic

deformation (20 to 40%). The a’ volume percentage is linearly re- lated to plastic deformation in Stage III. The nature of the transition in the roles of a' formation—from promoting earlier glide in

Stage II to being associated with a linear, high rate of work

30
hardening in Stage III —is not clear. Perhaps in Stage III essentially all active cross-slip sites have been transformed to a', and

subsequent a' formation occurs within active glide bands. Such a*

formation would not be expected to promote stacking-fault glide.

There is a hint that this sequence may occur: the x-ray data of Reed

and Guntner 56 presented schematically in Fig. 7 indicate that a ma- ximum value of the e volume concentration occurs at about the transition from Stage II to Stage III. There is apparently no e formation

during Stage III; the amount' of e decreases with strain, presumably

transforming to a

1 . 55 Stages I and II decrease in significance in

more stable austenites, but Stage III is retained. For instance,

AISI 316 exhibits neither Stage I nor Stage II, but does begin to

transform to a at strains of about 0.02 below 190 K. No e martensite

has been detected in this alloy.

In this discussion of the influence of martensitic products on plastic deformation, the results of current research indicate that

the schematic representations of Figs. 7 and 8 are time-dependent.

Further research is necessary to sort out the creep and strain-rate

effects and the isothermal characteristics of the strain- induced martensitic transformations.

This discussion has emphasized the roles of the strain-induced

martensites. Martensite that forms during cooling below T contributes differently. For many years a' martensite has been finown and

used to strengthen steel, and it may be used in a similar way to

strengthen Fe-Cr-Ni alloys (e.g., Mangonon and Thomas 52 ). With

severe rolling to induce a high percentage of a', room-temperature

yield strengths of about 1.4 GPa (2 x 10 5 psi) may be achieved at the

expense of significant ductility and (probably) toughness.

The complex martensitic transformation characteristics in

Fe-Cr-Ni austenitic commercial alloys are summarized in Table 2.

Stress-assisted martensite has been reported in single crystals 43 and

in Fe-9Cr-8Ni, 57 but has not been observed in Fe-18Cr-8Ni polycrystalline commercial grades. Martensite formed by the application

of elastic stress should result in morphology similar to that observed from cooling. The strain-induced transformations have been

subdivided into three stages, each contributing distinctly to austenite deformation. The major distinction between the stages is the

degree of work hardening. Strain-induced a* (Stage III) is distinguished from stress-assisted or thermally induced martensite by the

location of transformation: strain-induced a' forms at slip-band

intersections, with only one habit-plane variant; stress-assisted or thermally-induced martensite forms in {111} sheets with three habitplane variants. This review emphasizes one aspect that is not con- ventionally recognized. It has been thought that stress-assisted

transformation lowers the flow strength, and that strain-induced

transformation increases the flow strength. The strain-induced martensite at first promotes no work hardening and thus decreases the

31
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32

Stage

III

b.c.c.

Within

{111}

10-80 %

linear

with

strain;

sheets

strong

work

hardening
flow strength in Stage II, and then in Stage III, it increases the

flow strength. This leads one to suspect that the transformation

mechanism changes from Stage II to III.

INFLUENCE ON PERFORMANCE

Several characteristics of the austenite-to-b.c.c. -martensite

transformation in stainless steels cause problems in many cryogenic

applications

.

The martensitic b.c.c. product has a larger specific volume

(+1.7%) than the parent austenite; therefore, these steels expand

during transformation. The martensitic product forms as individual

crystals, and the associated transformation shear and volume expansion results in localized regions of disruption. For example, in

Fig. 9a, a* formation at the surface is illustrated microscopically.

The shape deformation results in local surface upheavals, usually

bounded by (111 )y planes. In Fig. 9b, a' formation near or at the

surface of the flange caused localized surface upheavals. For

service that requires close tolerances, such as valves or bearings,

these local surface fluctuations are disastrous.

Welding of austenitic steels tends to produce chemical segregation in the heat-affected zone. Nitrides, in high-nitrogen steels,

and carbides tend to form, especially with chromium. The chemically

depleted austenitic solid solution has a greater tendency for martensitic transformation on cooling to cryogenic temperatures. Transformation under such circumstances is particularly critical, since it is

restricted to the weld heat-affected zone and, therefore, tends to

produce high local stresses from the transformation volume expansion

and local shear. In a few applications of thin-gauge sheet, throughthickness cracks in the heat-affected zone of the welds or spot welds

have resulted.

The transformation at cryogenic temperatures is unpredictable

from room-temperature characteristics. The transformation occurs after application of applied stress or local plastic-deformation at

low temperatures much more readily than at room temperature. Prevention of in-service transformation depends on proper alloy selection. A good rule to follow is that the more alloying elements one uses (and can afford), the more stable the austenitic stainless steel

will be. Therefore, to a first approximation, AISI 310 is stable,

AISI 316 has limited strain-induced martensite, and AISI 304 is the

least stable. In AISI 304 within the chemical specification limits,

martensite may form on cooling and on application of stress and cer- tainly will form with plastic deformation.

Another concern, mostly in design of superconducting magnets, is

that the b.c.c. martensitic product is ferromagnetic, whereas the

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

33
34

Fig.

9.

(a)

Shape

deformation

of

martensitic

transformation

in

austenitic

stainless

steel

(Fe-Cr-Ni

alloy,

AISI

304L)

on

cooling

to

76K

(230X)

.

(b)

Shape

deformation

of

martensitic

transformation

in

steel

(AISI

304)

flange

previously

used

in

service

at

76K

(64X).
of b.c.c. martensite results in a permeability increase of 0.01.

Therefore, if time-dependent field changes corresponding to material

changes of the order of 0.10 are significant, then alloy selection is

important

.

Finally, a common design fear has been that martensitic transformation in austenitic stainless steels adversely affects toughness.

Experience suggests that this is partly true. Strain- induced martensitic transformation in austenitic stainless steels tends to produce higher toughness. Comparison of the low-temperature toughness

of AISI alloy grades 304, 316, 310 indicates that toughness is inversely related to austenite stability (and austenite tensile yield

strength). But, if the austenite is sufficiently unstable such that

martensite forms on cooling (e.g., in the AISI 400 series steels and

in AISI 301), then the martensite is less tough. To summarize: In

the more stable austenitic stainless steel grades, martensite does

not form during cooling and forms only during deformation. In these

alloys, with predominantly austenite plastic deformation, martensitic

transformation produces greater toughness. In metastable stainless

steels in which the base structure is predominantly martensite prior

to the application of stress and, therefore, in which plastic deformation must take place in the martensitic phase, the martensite has a deleterious effect on toughness.

REFERENCES

1. Cohen, M. , Olson, G. B., and Clapp, P. C., On the classification

of displacive phase transformations, in Proceedings of the

International Conference on Martensitic Transformations , ICOMAT 1979 , Department of Materials Science and Engineering,

M.I.T., Cambridge, Massachusetts (1979), pp. 1-11.

2. Reed, R. P. and Breedis, J. F. (1966). Low-Temperature phase

transformations, in Behavior of Materials at Cryogenic Temperatures , ASTM STP No. 387, American Society for Testing and

Materials, Philadelphia, Pennsylvania (1966), pp. 60-132.

3. Bilby, B. A. and Christian, J. W. , The crystallography of martensitic transformations, J. Iron Steel Inst. 197:122-131

(1961).

4. Christian, J. W. , Basic crystallography and kinetics, in

Martensite Fundamentals and Technology , E. R. Petty, ed.,

Longman Group, London, 11-41 (1970).

5. Wayman, C. M. , Introduction to the Crystallogrpahy of Martensitic Transformations , Macmillan Co., New York (1964).

6. Wayman, C. M. , The crystallography of martensitic transformations in alloys of iron, in Advances in Materials Research , Vol. 3, H. Herman, ed., Interscience, New York (1968),

pp. 147-304.

7. Lieberman, D. S., Crystal geometry and mechanisms of phase

transformations in crystalline solids, in Phase Transf orma35
tions , American Society for Metals, Metals Park, Ohio (1970),

pp. 1-58.

8. Roitburd, A. L. and Kurdjumov, G. V. , The nature of martensitic

transformations. Mater . Sci . Eng . 39:141-167 (1979).

9. Entwisle, A. R., The kinetics of martensite formation in steel,

Metall . Trans . 2:2395-2407 (1971).

10. Christian, J. W. , The Theory of Transformations in Metals and

Alloys , Pergamon Press, New York (1965).

11. Kaufman, L. and Cohen, M. , Thermodynamics and kinetics of

martensitic transformations, in Progress in Metal Physics , Vol. 7, B. Chalmers and R. King, eds., Pergamon Press,

New York, (1958) pp. 165-246.

12. Barrett, C. S., Transformations at low temperatures, Trans . Jap . Inst . Met . 17:465-475 (1976).

13. Martensite Fundamentals and Technology , E. R. Petty, ed.,

Longman Group, London (1970).

14. Christian, J. W. , The Theory of Transformations in Metals and

Alloys , Pergamon Press, New York (1976).

15. Physical Properties of Martensite and Bainite , The Iron and

Steel Institute Special Report 93, Percy Lund, Hampshires & Co., London (1965).

16. Mechanism of Phase Transformations in Crystalline Solids , Monograph and Report Series No. 33, Institute of Metals, London,

England (1969)

17. Shape Memory Effect in Alloys , J. Perkins, ed.. Plenum Press,

New York (1975)

18. New Aspects of Martensitic Transformations , H. Suzuki, ed.,

Japan Institute of Metals, Tokyo, Japan (1976).

19. Phase Transformations , American Society for Metals, Metals Park,

Ohio (1970).

20. Proceedings of the International Conference on Martensitic

Transformation , ICOMAT-1979 , W. S. Owen, ed.. Department of

Materials Science and Engineering, M.I.T., Cambridge, Massachusetts (1979)

21. Breedis, J. F. and Kaufman, L., Formation of hep and bcc phases

in austenitic iron alloys, Metall. Trans . 2:2359-2371 (1971).

22. Kaufman, L. , The free energy changes attending the martensitic

transformation in the iron-chromium and iron-chromium-nickel

systems, Trans AIME 215:218-223 (1959).

23. Kaufman, L. and Cohen, M. , The martensitic transformation in the

iron-nickel system, Trans . AIME 206:1393-1401 (1956).

24. Rao, M. M. , Russell, R. J., and Winchell, P. G., A correlation

of thermodynamic variables for iron-rich iron-nickel-carbon

alloys, Trans . AIME 239:634-642 (1967).

25. Kaufman, L., The lattice stability of the transition metals, in

Phase Stability in Metals and Alloys , P. S. Fudman,

J. Stringer, and R. I. Jaffee, eds., McGraw-Hill, New York

(1967), pp. 125-150.

26. Reed, R. P., The spontaneous martensitic transformations in

18%Cr, 8%Ni steels, Acta Metall. 10:865-877 (1962).

36
27. Eichelman, G. H. and Hull, F. C., The effect of composition on

the temperature of spontaneous transformation of austenite to

martensite in 18-8-type stainless steel, Trans . Am . Soc . Met . 45:77-104 (1953).

28. Monkman, F. C., Cuff, F. B. , and Grant N. J., Computation of M

for stainless steels. Met . Prog . 71:94-96 (1957).

29. Hammond, C. M. , The development of new high-strength stainless

steels, in Advances in the Technology of Stainless Steels and

Related Alloys , ASTM STP No. 369, American Society; for

Testing and Materials, Philadelphia, Pennsylvania (1963), pp.

47-53.

30. Andrews, K. W. , Empirical formulae for the calculation of some transformation temperatures, J. Iron Steel Inst. 203:721-727

(1965)

.

31. Hull, F. C., Delta ferrite and martensite formation in stainless

steels, Meld . J. Res . Suppl . 52 : 193-s-203-s (1973).

32. Angel, T., Formation of martensite in austenitic stainless

steels, J. Iron Steel Inst . 177:165-174 (1954).

33. Williams, I., Williams, R. G., and Capellaro, R. C., Stability

of austenitic stainless steels between 4 K and 373 K, in Proceedings of the Sixth International Cryogenic Engineering

Conference , IPC Science and Technology Press, Guildford,

Surrey, England (1976), pp. 337-341.

34. Warnes, L. A. A. and King, H. W. , The low temperature magnetic

properties of austenitic Fe-Cr-Ni alloys in the prediction of

Neel temperatures and maximum susceptibilities, Cryogenics

16:659-667 (1976).

35. Lee, S., Millman, S., MacDougall, I. L., Keown, P., and Argent,

B. B., Enthalpy of the martensitic transformation in steels

containing nickel and chromium, Met . Sci 11:261-271 (1977).

36. Dulieu, D. and Nutting, J., Stacking fault energy of Fe-Cr-Ni

alloys, in Metallurgical Developments in High Alloy Steels , Special Report 86, Iron and Steel Institute, London (1964),

pp. 140-145.

37. Schramm, R. E. and Reed, R. P., Stacking fault energies of seven commercial austenitic stainless steels, Metall. Trans.

6A: 1345-1351 (1975).

38. Rhode's, C. G. and Thompson, A. W. , The composition dependence of

stacking fault energy in austenitic stainless steels, Metall.

Trans . 8A: 1901-1906 (1977).

39. Brofman, P. J. and Ansell, G. S., On the effect of carbon on the

stacking fault energy of austenitic stainless steels, Metall.

Trans. A 9A:879-880 (1978).

40. King, H. W. and Larbalestier , D. C., Austenitic stainless steels

at cryogenic temperatures: The compositional dependence of

the M , Cryogenics 21:521-524 (1981).

41. Collings, E. W. and King, H. W. , Austenitic stainless steel as magnetic transition metal alloys, in The Metal Science of

Stainless Steels , E. W. Collings and H. W. King, eds.. Metallurgical Society of the AIME, New York (1979), pp. 1-21.

37
42. Bampton, C. C., Jones, L. P., and Loretto, M. H. , Stacking fault

energy measurements in some austenitic stainless steels, Acta

Metall . 26:39-51 (1978).

43. Breedis, J. F. , Martensitic transformations in iron-chromiumnickel alloys, Trans . AIME 230:1583-1596 (1964).

44. Kelly, P. M. , The martensite transformation in steels with low

stacking fault energy, Acta Metall . 13:635-646 (1965).

45. Wirth, A. and Bickerstaf f e , J., The morphology and substructure

of martensite in maraging steels, Metall. Trans. 5:799-808

(1974)

46. Lagneborg, R., The martensite transformation in 18%Cr-8%Ni

steels, Acta Metall . 12:823-843 (1964).

47. Abrassart, F. , Stress-induced y+a' martensitic transformation in

two carbon stainless steels. Application to trip steels,

Metall . Trans . 4:2205-2216 (1973).

48. Breedis, J. F. , Influence of dislocation substructure on the

martensitic transformation in stainless steel, Acta Metall . 13:2359-2371 (1965).

49. Strife, J. R., Carr, M. J., and Ansell, G. S., The effect of

austenite prestrain above the M temperature on the martensitic transformation in Fe-Ni-Cr-C alloys, Metall. Trans.

8A: 1471-1484 (1977).

50. Brooks, J. W., Loretto, M. H., and Smallman, R. E., Direct observations of martensite nuclei in stainless steel, Acta

Metall . 27 : 1839-1847 (1979).

51. Suzuki, T. , Kojima, H. , Suzuki, K., Hashimoto, T. , and Ichihara,

M. , An experimental study on the martensite nucleation and

growth in 18/8 stainless steel, Acta Metall . , 25, 1151-1162

(1977).

52. Mangonon, P. L. and Thomas, G., Struqture and properties of

thermal-mechanically treated 304 stainless steel, Metall

Trans . 1, 1587-1594 (1970).

53. Murr, L. E. , Staudhammer, K. P. and Hecker S. S., Effects of

strain state and strain rate on deformation-induced transformation in 304 stainless steel: Part II. Microstructural

study, Metall . Trans . 13A:627-635 (1982).

54. Mirzagev, D. A., Goykhanberg, Y. N. , Shteynberg, M. M. , and

Rushchin, S. V. , The elasto-plastic strain effects in alloys

with low stacking fault energy at temperatures above M , Fiz . Met . Metalloved . 35:1206-1212 (1973).

S

55. Guntner, C. J. and Reed, R. P., The effect of experimental variables including the martensitic transformation on the low

temperature mechanical properties of austenitic stainless

steels, ASM Trans . (). 55:399-419 (1962).

56. Reed, R. P. and Guntner, C. J., Stress-induced martensitic

transformations in 18Cr-8Ni steel, Trans. AIME 230:1713-1720

(1964).

57. Olson, G. B. and Azrin, M. , Transformation behavior of TRIP

steels, Metall . Trans

A 9A:713-721 (1978).

38
58. Breedis, J. F. and Robertson, W. D., Martensitic transformation

and plastic deformation in iron alloy single crystals, Acta

Metall. 11:547-559 (1963).

39

A COMPUTER- INTERACTIVE FATIGUE CRACK GROWTH RATE TEST PROCEDURE*+

Yi-Wen Cheng

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

A computer-interactive fatigue crack growth rate testing system has been developed that

implements the fatigue crack growth rate testing procedure set forth in ASTM E647-81. The system

uses a minicomputer to acquire the load-deflection data as a function of load cycles and to determine the crack length using the compliance technique. The fatigue crack growth rate, da/dN

and the stress intensity range, aK, are calculated and plotted during the test.

+Work sponsored by the Department of Interior, Minerals Management Serivce and Department of Energy,

Office of Magnetic Fusion Energy.

Key Words: compliance technique; computer- interactive testing; fatigue crack

growth rate; fatigue of materials.

41

INTRODUCTION

Fatigue crack growth rate (FCGR) data can be used for material characterization and for

fracture mechanics reliability analysis of structures subjected to cyclic loading. The standard

test method for measuring FCGR under constant-amplitude loading has been developed and published in

the 1981 Annual Book of ASTM Standards under the designation ASTM E647-81 [1]. In this paper, a computer-interactive implementation of the Standards is described.

There are several advantages of using a computer to acquire, analyze, and display FCGR data

during the test. First, the time needed for crack length determination and data reduction is greatly reduced. Second, the data scatter is reduced. Third, because the testing is interactive

in nature, the procedure is relatively easy to follow. Fourth, subjective interpretation and

influence of the experimenter are eliminated.

THE FCGR TEST METHOD

The requirements of specimen preparation for FCGR tests are described in ASTM E647-81 [1].

The sequence of the FCGR test is: first to obtain the raw data, namely, fatigue crack length, a, versus elapsed fatigue cycles, N; and then to analyze these data to a plot of da/dN versus aK,

I where da/dN is the FCGR in mm/cycle and aK is the crack-tip stress intensity range in MPa*m 2

. Typical outputs are presented in Figure 1.

The number of elapsed fatigue cycles is readily available from electronic or mechanical

counters. However, the methods of crack length measurement are more complicated and have been a subject of extensive study [2,3]. Although several methods of crack length measurement have been

developed, most require specialized equipment commonly not available in mechanical testing laboratories. The compliance technique, however, requires only monitoring of load cell and clip gage

deflection, which is routinely achieved in mechanical testing. Compliance is defined as the

specimen deflection per unit load, which is a function of crack length. The load and deflection

signals (voltages) can also be interfaced to a computer. Because of the simple instrumentation and

the need for environmental chambers for cryogenic temperature and saltwater corrosion-fatigue

tests, the compliance technique was chosen to measure the crack length.

The most commonly used FCGR test specimen is the compact- tension (CT) type with its wellestablished crack-tip stress intensity calibration [1] and compliance functions [4]. The data

reduction technique is either the secant method (point-to-point method) or the seven-point

incremental polynominal method, as recommended in ASTM E647-81 [1].

THE COMPUTER- INTERACTIVE FCGR TEST

Equipment and Procedure

A schematic of the computer-interactive FCGR testing system is illustrated in Figure 2. The

system consists of a closed-loop servohydraul ic mechanical testing machine, a multichannel

analog- to-digital (A/D) converter unit, and a minicomputer. Signal amplifiers and the cycle

43
counter are built into the mechanical testing machine. Included with the minicomputer are a cathode-ray-tube (CRT) terminal, a line printer, a dual floppy disk data storage unit, and a digital plotter.

The current A/D converter unit has a 16-bit resolution with a range of ± 10 V. This gives a least significant bit a resolution of about 0.3 mV. This level of precision has been found adequate

for the present purpose. At the present time, because of a 5-Hz filter associated with the A/D

converter unit, the maximum fatigue cycle frequency that the A/D converter can convert enough loaddeflection pairs for an accurate crack length measurement is limited to about 0.03 Hz. Consequently, fully automatic testing with higher frequencies is not yet possible, and the test must be periodically interrupted for crack length measurement. A digital memory oscilloscope has

been procured and is being incorporated into the testing system. It will enable the testing system

to be fully automatic during continuous high-frequency testing.

The procedure for conducting the FC6R test is described as follows. The precracked specimen

is fatigue cycled under the prescribed loading conditions and frequency. A typical cycling

frequency is 20 Hz. Depending upon the crack length-to-specimen width ratio, a/W, the fatigue

cycling is interrupted after selected increments of crack growth. The instantaneous crack length

is then inferred using the compliance technique. For example, in a 25. 4 -mm- thick CT specimen,

crack length measurements at an interval of crack growth between 0.5 mm and 0.76 mm are optimum.

The amount of crack length increase during the 20-Hz fatigue cycling can be estimated using a digital indicator, which provides continuous instantaneous readings of the maximum and minimum

output voltages from the clip gage at the crack mouth. The difference in maximum and minimum

values correlates well with the instantaneous crack length.

During the single 0.02-Hz cycle compliance measurement, the signals of load cell and clip gage

are plotted on an X-Y recorder. At the same time, the signals are amplified, digitized by the A/D

converter unit, and fed into the computer for compliance calculation. The lower one- third of the

compliance signals are dropped to eliminate the possible crack closure effects [5]. The upper two- thirds of the signals are correlated to a straight line using a linear least-squares fit. This is illustrated schematically in Figure 3. A linear correlation coefficient of 0.9999 or greater is generally obtainable. If the linear correlation coefficient drops below 0.999 scatter will

increase so the test is terminated. Generally, a drop in linear correlation is obtained at high aK

values in ductile steels because of the development of a large plastic zone ahead of the crack tip.

The large plastic zone causes the specimen to behave inelastically resulting in the formation of a hysteresis loop in the load-deflection curve. From the resulting compliance, the instantaneous

crack length is computed using the appropriate expression for the compliance cal ibration .of the

specimen [4]. The precision of the crack length measurement is typically within ± 0.04 mm.

The specimen compliance, the least-squares correlation coefficient, and the crack length

inference are displayed on the CRT screen and printed on the line printer. After the number of

fatigue cycles is manually fed into the computer through the CRT terminal, the computer computes

da/dN and aK, and the digital plotter plots the data points (a,N) and (da/dN, aK) on the a-versus-N

and on the da/dN- versus-AK graphs, such as those in Figure 1. The line printer also prints the

da/dN and aK results. All the resulting data are stored on the floppy disk for post-test analyses.

44
The computer programs for the post-test analysis include the following capabilities:

1. Converting units

2. Plotting data in desired units

3. Plotting data in desired coordinate ranges

4. Plotting data of several different specimens on one graph

5. Calculating the material constants C and m in the Paris equation, da/dN = C (aK)

171 [6],

and drawing the regression line through the data

All the computer programs, including data acquisition routine, data reduction routine, and the

post-test analysis routines, are written in the language of Fortran IV.

APPLICATIONS

The computer- interactive FCGR testing system has been successfully used for room temperature

tests in air, cryogenic temperature tests in liquid nitrogen and in liquid helium, and room temperature tests in saltwater. Special apparatus are needed in the cryogenic temperature tests

and in the saltwater tests. The loading apparatus used in the cryogenic temperature tests have

been described previously by Fowlkes and Tobler [7]. In the saltwater tests, the clip gage used

for deflection measurements is mounted above the saltwater level on a scissors-like extension that

protects the gage from the saltwater.

Figure 4 shows a comparison between the results obtained with the computer- interactive FCGR

testing system and those obtained without the computer system, using manual data reduction. Good

agreement between the two tests is observed, but the computer-interactive testing accomplished the

tasks in far less time.

SUMMARY

A computer- interactive FCGR testing system, which implements the procedure set forth in ASTM

E647-81, has been developed and successfully used in room temperature tests in air, in cryogenic

temperature tests, and in room temperature tests in saltwater. The system offers considerable time

saving in data acquisition and in data reduction. The new test procedure is relatively easy to follow and enables technicians to produce data with less scatter (with respect to the

non-computer-aided technique), because no manual data interpretation or fitting to straight lines

is done.

45
REFERENCES

[1] Standard Test Method for Constant-Load-Ampl i tude Fatigue Crack Growth Rates Above 10" m/cycle, (Designation E647-81) 1981 Annual Book of ASTM Standards, Part 10, American Society

for Testing and Materials, Philadelphia, PA, (1981), pp. 765-783.

[2] The Measurement of Crack Length and Shape During Fracture and Fatigue, C. J. Beevers, ed.

Engineering Materials Advisory Services Ltd. Printed by the Chameleon Press, Ltd, UK (1980).

[3] Advances in Crack Length Measurement, C. J. Beevers, ed. Engineering Materials Advisory

Services Ltd., Printed by the Chameleon Press, Ltd, UK (1981).

[4] S. J. Hudak, Jr., A Saxena, R. J. Bucci, and R. C. Malcolm, "Development of Standard Methods

of Testing and Analyzing Fatigue Crack Growth Rate Data," Air Force Materials Laboratory,

AFML-TR-78-40, (May, 1978).

[5] W. Elber, "The Significance of Fatigue Crack Closure," in Damage Tolerance in Aircraft

Structures , ASTM STP 481, American Society for Testing and Materials, (1971), pp. 230-242.

[6] P. C. Paris and F. Erdogan, "A Critical Analysis of Crack Propagation Laws," Transactions of

ASME, Journal of Basic Engineering , Series D, Vol . 85, No. 3 (1963), pp. 528-534.

[7] C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a Ti-6al-4V Alloy at Liquid

Helium Temperature," Engineering Fracture Mechanics , 8 (1976), pp. 487-500.

46
LIST OF FIGURES

1. Data outputs from the FCGR test.

2. Schematic of the FCGR testing system.

3. Schematic illustration of the closure effect on the load-deflection

curve.

4. Results of FCGR test obtained from the computer-interactive testing system and the

non-computer-aided technique: a comparison.

47
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49
DEFLECTION

Fig. 3. Schematic illustration of the closure effect on the loaddeflection curve.

50
da/dN,

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13Cr-19Mn Stainless Steel -

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51

EFFECTS OF SENSITIZATION ON SOME FRACTURE MECHANICS FATIGUE

PARAMETERS FOR Fe-18Cr-3Ni-13Mn-0.37N STAINLESS STEEL

AT 295 AND 4 K*

R. L. Tobler

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

A nitrogen-bearing austenitic stainless steel (Fe-18Cr-3Ni-13Mn-0.37N) was subjected to fatigue tests at 295 and 4 K. Sensitization (chromium carbo- nitride precipitation) was induced by

heating the as-received annealed material for 0.75 or 1.5 h at 977 K. Data comparisons for the

sensitized versus the annealed conditions indicate that at 295 K there is little or no difference

in the Paris law parameters governing fatigue crack growth rates at intermediate aK. At 4 K, however, sensitization reduces the fatigue resistance of this steel by adversely affecting all

three stages of the fatigue process: crack initiation, crack propagation, and fatigue fracture

toughness.

Key words: austenitic stainless steel; chromium carbo-nitride precipitation; fatigue;

fatigue crack growth rates; low temperature properties; sensitization.

*Work supported by Office of Fusion Energy, Department of Energy.

53

INTRODUCTION

The austenitic stainless steels are generally regarded as reliable cryogenic structural alloys, but property degradation due to sensitization may limit the performance of some of these

materials for specific applications, such as those requiring welding. Sensitization is a metallurgical condition encountered when austenitic stainless steels are heated at temperatures from 700

to 1100 K. After sufficient time at these temperatures, chromium carbo-ni tride compounds begin to precipitate, first at the grain boundaries, and then within the austenite grains themselves. This

precipitation can ruin corrosion resistance, reduce ductility and toughness, and lower the austenite stability with respect to martensite transformations.

Little is known about the effects of sensitization on the fatigue properties of austenitic

stainless steels at cryogenic temperatures. Some steels are more susceptible to sensitization,

owing to compositional variations. In this paper some measurements are reported for an Fe-18Cr3Ni-13Mn-0.37N stainless steel. This relatively high-nitrogen alloyed steel was studied because it was thought to be particularly susceptible to sensitization at 4 K.

MATERIAL

The test material was a commercial 25-mm plate of austenitic stainless steel, previously

tested in the annealed condition and characterized [1] as follows:

As- received condition : annealed, 1080°C, water quenched

Composition : (wt%) Fe-18.09Cr-3.26Ni-13.22Mn-0.37N-0.038C-0. 12Mo-0.52Si-0.005S-0.028P

Grain Size : ASTM No. 6 (d = 48 ym)

Hardness : Rockwell B 93

Mechanical properties : See Table 1

This plate was tested in both the annealed (as-received and not sensitized) and sensitized

conditions. To sensitize some of the material, separate portions of stock were cut into blanks and

then subjected to either of two thermal treatments:

S^ condition: 977 K, 0.75 h, water quench

S

2 condition: 977 K, 1.5 h, water quench

Thus, the annealed condition is unsensitized, while the S^ condition is more severely sensitized than the condition. Figure 1 shows the microstructures as revealed by electrolytic

etching with 10 % oxalic acid solution for various periods of time. In addition to the usual

examination with the light microscope, specimen was heavily etched and examined in the scanning

electron microscope (SEM).

55
PROCEDURE

All tests were conducted using compact specimens 25-mm thick and 50.8-mm wide, as shown in Figure 2. The specimen orientation is such that the load is applied in the transverse direction

and the crack propagates in the longitudinal direction of the plate. The specimen, clip-gage, and

cryogenic apparatus for performing tests in liquid helium at 4 K are described elsewhere [2].

The properties measured in this study included the number of cycles to initiate a fatigue

crack from the notch (N.), the fatigue crack growth rates (da/dN), and the fatigue fracture tough- ness (Kf ). These parameters were measured at selected temperatures and test conditions specified

in the text. After failure the specimens were examined with the SEM.

All fatigue tests used a sinusoidal wave form and a 0.1 ratio of minimum- to-maximun fatigue

loads. The cycling rate was 20 Hz, except near the end of the tests when higher stress intensity

factor (K) levels were approached. Then, the frequency was reduced to values as low as 2 Hz.

The fatigue crack initiation life was determined for two annealed and two sensitized specimens

at 4 K by recording the compl iance-versus-cycles data and reading the N. values from the resulting

curves. A marked change in compliance (deflection/load) occurs when a crack develops in this

material at 4 K, leading to relatively easy identification of the crack initiation life at this

temperature. The notch root radii were 0.127 mm ± 10 %. The nominal AK value at initiation was calculated using the starting notch length as the crack length.

Subsequent to fatigue crack initiation, the crack growth rates were measured at 295 or 4 K, using Cheng's computer-aided compliance test method [3]. The expression used to calculate the

stress intensity factors is given in reference [4]. The resulting da/dN - versus - AK data were

plotted on log-log coordinates and the Paris equation parameters C and n (da/dN = C AK

n

) were determined on the basis of least-squares fits.

Examination of the specimens after testing showed that obtaining uniform crack fronts in this

25-mm thick material at 4 K is a problem. The ratios of edge-to-average crack lengths at 4 K were

always less than 0.9. However, it is assumed that any effects of crack front curvature are satisfactorily accounted for in the fatigue crack growth rate tests since the compliance method senses

an average crack length.

The fatigue fracture toughness [5], , was estimated from observations made at the termination of da/dN tests. Fatigue cycling with periodic compliance readings was continued until frac- ture occurred, at which time the final cycle number was recorded. The crack length at failure was estimated from the final cycle number and the da/dN- versus-Ak trend for the specimen. Calculations

of Kp were then based on the maximum fatigue load and the estimated crack length at failure. The

uncertainty in K^

c due to error in the final average crack length is estimated at ± 10 °%.

Fatigue fracture toughness tests were not performed at 295 K because this steel is highly ductile at room temperature. At 4 K this steel exhibits linear-elastic behavior, and the thickness of

25 mm is quite sufficient to satisfy the ASTM E- 399- 81 [6] size requirements for valid fracture

56
toughness (Kj ) tests. (The thickness requirement for the annealed condition at 4 K is only

5.3 mm). Although the amount of crack front curvature at 4 K is sufficient to invalidate Kj

c mea- surements [6], the Kp measurements of this study are still considered useful as qualitative indicators of the effect of sensitization on toughness.

RESULTS

Fatigue Crack Initiation: The fatigue crack initiation data at 4 K are shown in Figure 3.

i Tests were conducted for only one nominal AK value of 20.5 MPa*m 2

, calculated using the initial

notch length as the crack length. For two specimens in the annealed (unsensitized) condition, the

initiation lives were 95 and 127 kilocycles (kc). Single tests for the sensitized specimens gave

N. = 85 kc for the S-^ condition and = 36 kc for the S£ condition. The results show that fewer

cycles are required to initiate a fatigue crack when the material is sensitized. This effect is

i observed at aK = 20.5 MPa*m 2

, but the behavior at other stress intensity factor ranges remains

unexpl ored.

Fatigue Crack Growth : Fatigue crack growth rates for the two sensitized conditions at 295 and

4 K are shown in Figure 4. The Paris equation line fits for the S-^ and S

2 conditions at 295 K are essentially equivalent and in agreement with the results for annealed material from the same plate

of steel [1]. One explanation for the similar fatigue crack growth rate behavior at room temperature is the constancy of the failure mechanism. Figure 5 is a representative illustration of the

transgranul ar failure mode observed for each condition at room temperature, the distinguishing

feature being striation formation. This same failure mode is observed for the steel in all three

metallurgical conditions: annealed, and S

2

. Since this failure mode does not involve crack

propagation along grain boundaries, the effect of grain boundary precipitation on the room temperature cracking rates appears to be negligible for the aK range studied.

The temperature reduction from 295 to 4 K greatly accelerates the fatigue crack growth rates

of the Fe-18Cr-3Ni-13Mn-0.37N steel. This is shown clearly for the sensitized conditions (both

and S

2

) in Figure 4, and it is also reported to be the case for the annealed condition [1]. The

change from 295 to 4 K has remarkable effects on the Paris equation parameters also, in that C decreases while n increases markedly (Table 2) for both the annealed and sensitized conditions.

Changes in the failure modes associated with the temperature reduction are documented in Fig. 6: the ductile striation formation observed at 295 K is replaced at 4 K by more brittle mechanisms

which cause higher growth rates. The brittle failure modes at 4 K are distinguished by the

following features: 1) cleavage- type facets in the case of the annealed material (Fig. 6A), and

2) varying amounts of intergranular facets in the case of the sensitized material, conditions

(Fig. 6B) and S

2 (Fig. 6C). Thus, the relative brittleness of the fatigue failure mechanisms at 4

K explain the higher growth rates at 4 K, as compared to 295 K.

The fatigue crack growth rates at 4 K for the annealed and sensitized conditions are compared

in Figure 7. The data for annealed material derive from tests performed on three specimens,

whereas the data for the and S

2 conditions represent single specimen tests. For the aK range

studied, sensitization tends to increase the da/dN at 4 K for a given value of AK. Most of the

data for the sensitized conditions lie above the data scatter for the three annealed specimens.

The difference in rates varies with aK, as well as the degree of sensitization. In particular, the

57
rates for the S

2 condition begin to accelerate well above the rates for both the and annealed

conditions when aK exceeds 30 MPa*m 2

. As indicated below, the 4 K fatigue fracture toughness of

the sensitized material in the condition is relatively low (K

fc = 41 MPa*m 2 ), and this factor

apparently promotes the marked acceleration of da/dN for aK above 30 MPa*m 2

.

A summary of the da/dN- versus-AK results is given in Table 2. The regression statistics

listed in Table 2 apply to the Paris equation line fits which are illustrated in Figures 4 and 7.

Fatigue Fracture Toughness : The fatigue fracture toughness estimates, K^, indicate that

sensitization can significantly decrease the toughness, relative to the annealed condition, if

sufficient time (1.5 h) at the sensitizing temperature is allowed. The K

f values are 61, 61, and

I TC

41 MPa*m 2 for the annealed, S^, and conditions, respectively.

Decreased toughness in the sensitized condition was expected, since precipitates in the grain

boundaries usually lower the boundary cohesive strength. The results of scanning electron microscopy, Fig. 8, bear this out. These fractographs correspond to the overload failure zones for each

specimen condition. In each case the 4 K fracture surface consists of brittle failure modes but

there are obvious differences in the fracture features depending on material condition:

1) The annealed condition, Fig. 8A, shows a mixed mode consisting of cleavage facets interspersed with dimpled zones. Slip line traces are observable on the facets.

2) The sensitized condition. Fig. 8B, also shows a mixed mode, this time consisting of

portions of intergranular failure interspersed with dimpled zones.

3) The more severely sensitized condition. Fig. 8C, exhibits the most brittle behavior of

all. Here the intergranular failure is fully developed, predominating over essentially

100 % of the fracture surface. Many of the exposed grains evidence cavities on their

boundaries where the carbo-nitride particles had precipitated.

Thus, the effect of sensitization is to lower the fatigue fracture toughness at 4 K by pro- moting intergranular fracture in this steel.

DISCUSSION

The purpose of this discussion is to compare the fatigue and static fracture toughnesses of

this austenitic steel at 4 K. Sometimes it is assumed that failure will occur when K in fatigue

Ilia X equals K^. The present data indicate that this is not true for the studied material.

1 The K^

c value for the annealed specimen at 4 K was 61 MPa*m 2

. For the same heat of Fe-18Cr3N i - 1 3Mn- 0 . 37N steel, material condition, specimen thickness, and specimen orientation (TL[6]),

1 three static fracture toughness tests gave Kq values of 70, 70, and 73 MPa*m 2 [1]. These are Kq, not Kj measurements, owing to excessive crack front curvature [1,6]. On the other hand, the

load-deflection curves are linear-elastic and the thickness criterion is amply satisfied, so that

58
the Kg data do provide qualitative evidence that the static fracture toughness is higher than the

fatigue fracture toughness. Apparently then, K fc

s 0.85 Kg for this steel at 4 K.

A plastic zone size correction (r^) to the final crack length and the fatigue fracture tough

ness value can be made, based on the expression [7]:

This correction raises the K, estimate for the annealed condition at 4 K from 61 to 62

i f c MPa*m 2

. Since such a correction is not sufficient to account for the difference in toughness

between K^ and Kg, it appears from the data for the annealed condition that the static fracture

toughness is definitely higher (^16 %) than the fracture toughness that is measured during fatigue

as the crack grows to critical size at progressively increasing aK values.

SUMMARY AND CONCLUSION

Some effects of sensitization on the fatigue resistance of an austenitic stainless steel at liquid helium temperature were explored using 25-mm thick compact specimens. The test material was nitrogen-bearing Fe-18Cr-3Ni-13Mn-0.37N steel. Fatigue crack growth rates and Paris equation

parameters are reported at 295 and 4 K, along with a limited number of fatigue crack initiation and

fatigue fracture toughness measurements at 4 K.

1 For the aK range 20 to 50 MPa*m 2

, sensitization has little or no effect on fatigue crack

growth rates at room temperature. At 4 K, however, the rates for sensitized and annealed

conditions increase drastically compared to the rates at 295 K, and the effects of sensitization

relative to the annealed condition are then quite pronounced. Observations for the

Fe- 18Cr-3Ni- 13Mn-0.37N steel at 4 K are sufficient to demonstrate that sensitization (1.5 h at

977 K) leads to:

1) reduced crack initiation resistance (^60 %)

2) reduced crack growth resistance (a factor of 2 or more), and

3) reduced fatigue fracture toughness (^33 %).

Fractography demonstrates that the grain boundary precipitation resulting from sensitization has

little effect on the fatigue behavior at 295 K since the fatigue failures at this temperature occur

via a ductile transgranul ar striation mechanism. During fatigue and fracture testing at 4 K, however, the annealed material exhibits a brittle transgranular failure mode involving cleavage;

the effect of sensitization at this temperature is to promote a transition to a still more brittle

intergranular failure mode.

In conclusion, a decreased fatigue life at cryogenic temperatures may be anticipated for

sensitized Fe-18Cr-3Ni-13Mn-0.37N stainless steel, owing to a combination of effects. Additional

data would be necessary to fully quantify the effects of composition, test temperature, severity of

sensitization, and magnitude of fatigue stresses.

59
ACKNOWLEDGEMENT

The assistance of Mr. P. Purtscher of NBS in preparation of metallographic specimens and

fractography is gratefully acknowledged.

REFERENCES

[1] Read, D. T. and Reed, R. P., Toughness, Fatigue Crack Growth, and Tensile Properties of Three

Nitrogen-Strengthened Stainless Steels at Cryogenic Temperatures , in: The Metal Science of

Stainless Steels, E. W. Collings and H. W. King, eds., Met. Soc. AIME, 1979, pp. 92-121.

[2] Read, D. T. and Tobler, R. L., Mechanical Property Measurements at Low Temperatures, in: Adv.

Cryo. Eng. 28, R. P. Reed and A. F. Clark, eds., Plenum 1982, pp. 17-28.

[3] Cheng, V. W., A Computer- Interact!' ve Fatigue Crack Growth Rate Test, in this NBSIR volume.

[4] Hudak, S. J., Jr., Saxena, A., Bucci, R. J., and Malcolm, R. C., Development of Standard

Methods of Testing and Analyzing Fatigue Crack Growth Rate Data, Tech. Report AFML-TR-78-40,

May 1978, pp. 257, ADA058320.

[5] Yokobori, T. and Aizawa, T., Rep. Res. Inst. Strength Fract. Mater., Tohoku University, 6_, 1970, p. 19.

[6] Standard Test Method for Plane-Strain Fracture Toughness of Metallic Materials, ASTM Designation E 399-81, 1981 Annual Book of ASTM Standards, Amer. Soc. Test. Mater., Philadelphia,

1981, pp. 588-618.

[7] Irwin, G. R., Plastic zone near a crack and fracture toughness, in: Proceedings, 7th

Sagamore Ordinance Materials Research Conference, Syracuse University Press, Syracuse, NY,

1960, pp. iv63-iv78.

60
LIST OF TABLES

1 .

2 .

Mechanical properties for annealed

Fatigue crack growth rate data for

Fe-18Cr-3Ni-13Mn-0.37N stainless steel

Fe- 18Cr- 13Mn-0. 37N stainless steel.

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LIST OF FIGURES

1. Microstructures of annealed and sensitized Fe- 18Cr-3Ni- 13Mn-0. 37N steel: (A) annealed, 400X;

(B) S„ condition, 150X; (C) S^ condition, 200X; (D) S. condition, 600X; (E) S, condition, SEM

view after extensive etching, 500X; (F) same as E; pits have formed where the precipi tates

have been etched out.

2. Compact specimen used in this study.

3. Fatigue crack initiation test results, showing reduced life for the sensitized conditions.

4. Fatigue crack growth rate data for sensitized conditions at 295 and 4 K.

5. Fatigue failure surface at 295 K, showing tlje ductile transgranul ar striation mechanism for

material in the condition (aK s 50 MPa*m 2 ), (A) 300X; (B) 1000X.

6. Fatigue failure surfaces A at 4 K, showing mixed modes with brittlf mechanisms: (A) annealed

condition, x aK s 35 MPa*m 2

, 500X; (B) S, condition, aK s 25 MPa*m 2

, 300X; (C) S

? condition, aK. s 30 MPa-m5

, 750X.

1 c

7. Fatigue crack growth rate data at 4 K.

8. Fatigue surfaces corresponding to overload at 4 K: (A) annealed condition, 500X; (B) S, condition, 500 X; (C) S£ condition, 1000X.

63
Figure 1. Microstructures of annealed and sensitized Fe-18Cr-3Ni-13Mn-0. 37N

steel: (A) annealed, 400X; (B) S

? condition, 150X; (C) S, condition, 200X; (D) S, condition, 600X; (E) S, condition, SEMview

after extensive etching, 500X; (F) same as E; pits have formed

where the precipitates have been etched out.

64
COMPACT SPECIMEN

Thickness = 24.5 mm

Figure 2. Compact specimen used in this study.

65
COMPLIANCE,

S/P,

mm/kN

Figure 3. Fatigue crack initiation test results, showing reduced life for the

sensitized conditions.

66
FATIGUE

CRACK

GROWTH

RATE,

da/dN,

mm/cycle

STRESS INTENSITY FACTOR RANGE, AK, MPa-m1/2

_2 10 50

Figure 4. Fatigue crack growth rate data for sensitized conditions at 295 and

4 K.

67

FATIGUE

CRACK

GROWTH

RATE,

da/dN,

in/cycle
Figure 5. Fatigue failure surface at 295 K, showing the ductile transgranul ar stria^ion mechanism for material in the S, condition ( a

= 50

MPa-m 2 ), (A) 300X; (B) 1000X. 1

68
A

60ftm

Figure 6. Fatigue failure surfaces at 4 K, showing mixed modes with brittle

mechanisms: (A) annealedcCondi tion , aK = 35 MPa*m°, 500X; (B) S ia

condition, aK = 25 MPa*m , 300X; (C) S

? condition, aK = 30 MPa*m

750X

.

69
FATIGUE

CRACK

GROWTH

RATE,

da/dN,

mm/cycle

STRESS INTENSITY FACTOR RANGE, AK, ksMn1/2

-2 10 50

STRESS INTENSITY FACTOR RANGE, AK, MPa-m1/2

Figure 7. Fatigue crack growth rate data at 4 K.

70

FATIGUE

CRACK

GROWTH

RATE,

da/dN,

in/cycle
Figure 8. Fatigue surfaces corresponding to overload at 4 K: (A) annealed

condition, 500X; (B) condition, 500 X; (C) condition, 1000X.

71

EFFECTS OF MAGNETIC FIELD ON TENSILE BEHAVIOR AT 4 K

OF ALLOYS 304 AND 310+*

R. P. Reed, J. M. Arvidson, J. W. Ekin, and R. H. Schoon

National Bureau of Standards

Boulder, Colorado

ABSTRACT

The application of magnetic fields has been demonstrated to enhance

face-centered cubic to body-centered cubic (bcc) martensitic transformations

in ferrous austenitic alloys. In Fe-Ni-Mn alloys higher fields serve to raise

the temperature of possible transformation and to increase the amount of

transformation. There has been an abundance of papers discussing the use of

strong (up to 35 T) pulsed fields to enhance transformation in Fe-Ni, Fe-Ni-Cr

and Fe-Mn steels. The plastic deformation of a series of Fe-18Cr, 8- 25N

i

alloys at 4 K exposed to steady-state and transient magnetic fields has been

previously studied and small changes of flow strength were observed that have

been attributed to magnetostriction and joule heating effects. Many experiments have shown that martensitic transformations in Fe-Cr-Ni stainless steels

affect the low temperature flow strength.

Our experiments were conducted to assess the effects of a steady magnetic

field on the austenite stress-strain characteristics at 4 K. Wire specimens

(0.64-mm dia) of both a stable Fe-26Cr-20Ni (AISI 310) and metastabTe Fe-18Cr9Ni (AISI 304) were measured. A 7 T magnetic field was produced by a splitpair NbTi superconducting magnet with radial access ports and the field was applied perpendicular to the tensile specimen axis. Tests were conducted at

4 K with the field off, with the field on, and with the field switched on and

off after prescribed amounts of plastic deformation. The Young's modulus,

tensile yield strength and tensile stress-strain curves were measured. The

results are summarized here: (1) During magnetic field on/off experiments, no change in austenite flow strength of either alloy was observed from the

application or absence of a 7-T field and (2) There was no detectable effect

of a constant 7-T magnetic field on the yield strength and Young's modulus of

either alloy. However, the specimen-to-specimen yield strength data scatter

was large and the application of magnetic field would not necessarily produce

changes of flow strength in excess of the data scatter. Discussion is also

presented on magnetic field effects on discontinuous yielding and cold worked

alloys.

+ Work supported by Office of Fusion Energy, Department of Energy

* To be published in Austenitic Steels at Low Temperatures, Eds. R. P. Reed

and T. Horiuchi, Plenum Press, New York (1983).

73

INTRODUCTION

Superconducting magnet structures require alloys that have

high strength at 4 K in magnetic fields. Since magnetic fields

are known to influence martensitic transformation in steels, it is

prudent to examine the effect of magnetic field on their strength.

It has been demonstrated that magnetic fields tend to raise

the martensitic transformation temperature and to increase tha

amount of transformation in ferrous alloys. Korenko and Cohen v have recently studied an Fe-22 . 5Ni-4Mn alloy and showed that the

magnitude of the applied field significantly influences transformation kinetics. For example, no martensite was detected in this

alloy on cooling to 190 K; however, a constant 8-T field, applied

after cooling, produced 5 percent martensite, and a constant 14-T

field produced 35 percent martensite at 190 K.

/ o_c ) There has been an abundance of papers' discussing the use

of strong (up to 50-T) pulsed fields to enhance transformation in

Fe-Ni, Fe-Ni-Cr, and Fe-Mn steels. These studies have established

that pulsed magnetic fields raise the temperature of martensitic

transformation and, for a given temperature and alloy, there is a minimum critical field to produce martensite. The application of

magnetic fields influences only austenite to body-centered cubic

(ferromagnetic) martensite, not austenite to hexagonal close-packed

(nonferromagnetic) martensite.

*This research was supported by the Office of Fusion Energy, U.S.

Department of Energy.

75
The studies suggest that the primary factor for the influence of

magnetic field on transformation is that the thermodynamic stability between the alloy phases has been altered.

(9) Bolshutkin, Desnenko, and Ilichev have studied the plastic

deformation of a series of Fe-18Cr, 8 to 25Ni alloys at 4 K exposed

to steady-state and pulsed magnetic fields. No change of flow

strength was observed within experimental uncertainty (3-4 percent)

on application of a constant 3.4-T magnetic field. However, when

the magnetic field was applied during straining at 4 K, the flow

strength of Fe-18Cr, 8 to 15Ni alloys was observed to decrease.

The amount of strength decrease was between 1.0 and 5.0 MPa,

reaching about 0.6 percent of the flow strength. This decrease of

flow strength only became detectable after plastic strains of

about 0.05 and increased rapidly with increasing strain to a strain of about 0.10, where the magnitude of the decrease became

constant. Conversely, they found that in more stable alloys (Fe,

18Cr, 20-25Ni) the flow strength increased with application of a

3.4-

T field during testing. In these steels, the increase of flow

strength varied from 2.0 to 3.0 MPa, beginning at strains near

0.01 and saturating at strains near 0.02.

The effect of both longitudinal and transverse, constant

2.5-

T magnetic fields on the plastic deformation of high-purity

aluminum at 4 K has been studied. Increases of flow strength

during straining, associated with the application of the magnetic

field, were observed. The flow strength increase was associated

with increased electron deceleration, which results in reduced

dislocation velocity.

The effect of magnetic field on the flow strength at^.K has

been recently studied for a ferromagnetic metal, nickel.

Magnetic fields were found to increase the flow strength, application of a transverse field had greater influence than a parallel

field, and the flow strength increased linearly with field in the

range 0 to 4 T. Resistivity measurements as a function of strain

provided evidence that the dislocation structure was not affected

by magnetic field. The increased flow strength under magnetic

fields is attributed to the generation of spin waves by moving

dislocations in the ferromagnetic lattice, resulting in reduced

dislocation velocity.

This study was initiated to begin assessment of the effects

of magnetic fields on austenite-martensite deformation at 4 K in

Fe-Cr-Ni alloys. Many experiments have shown that strain-induced

martensitic transformations in Fe-Cr-Ni stainless steels affect

the low temperature flow strength. At low strains, transformation

apparently results in reduced flow strength. However, at larger

strains, bcc martensite transformation correlates excellently with

increased work hardening. Also, elastic property and

76
susceptibility measurements have indicated the complex magnetic

state of Fe-Cr-Ni austenites at low temperatures. This state,

below the Ndel temperature, has been described as varying from

longer-range antiferromagnetic ordering in Fe-18Cr-8Ni alloys to

shorter-range spin-glass interactions in alloys containing more Cr

or Ni. Therefore, two major questions need to be answered: (1) Do

applied fields enhance strain-induced martensite transformation in

Fe-18Cr-8Ni steels and, if so, how does the magnetic-field-enhanced

martensite affect austenite flow strength? (2) How do applied magnetic fields affect the flow strength of austenite as a function of

composition?

MATERIALS

Two stainless steels, AISI 304 (Fe-19Cr-8Ni) and AISI 310

(Fe-25Cr-20Ni) were selected for this test series. AISI 304 is

metastable with respect to austenite-to-martensite transformation

and transforms during plastic deformation at temperatures below

about 200 K; AISI 310 is completely stable with respect to martensitic transformation. Both were obtained as 0.64 mm diameter wire

in the cold-drawn condition. Annealing times of 7.5, 15, and

30 min at 1340 K (1950 °F) produced yield strength reductions from

400 to 305, 285, and 290 MPa, respectively, at room temperature

for alloy 310. Annealing of alloy 304 at 1340 K (1950 °F) for

30 min resulted in a yield strength reduction from 321 to 233 MPa

at room temperature. All annealed wire specimens were produced by

heating in a vacuum at 1340 K (1950 °F) for 30 min.

EXPERIMENTAL PROCEDURE

(14) The apparatus has been described previously by Ekin.

Tensile load and a magnetic field were simultaneously applied to

wire specimens at 4 K. The magnetic field was produced by a 7-T

split-pair NbTi magnet with radial access ports, allowing perpendicular application of the field to the specimen. A schematic of

the apparatus is shown in Fig. 1. Specimens were tested in

straight sections, each with a 27.5-cm-long gage length. The

profile of the magnetic field strength over the specimen length is

illustrated in Fig. 1. The magnet produces a 7-T field perpendicular to a 6.3-cm length of the specimen at its center. On each

side of this central section, the field strength decreased linearly

over an 8 cm length to 0 T. Therefore, while only about h of the

specimen is exposed to 7 T, the entire specimen is exposed to some magnetic field. Magnetic field accuracy was ±1.0 percent, and

homogeneity over the sample central section was of the order of

± 1.0 percent.

77
—B(T)

7 0

I 1

Magnetic

Field Profile Wire Specimen

t

\ / Split Pair

V Superconducting

/ \ Magnet

Load

Fig. 1. Schematic of apparatus to apply tensile force and magnetic

field to wire specimens.

The specimens were gripped at both ends by soft silver soldering

into copper channels, one fixed and one attached to the load train

through a self-aligning universal joint. Strain was measured by

taking the vertical displacement of the pull rod, as indicated by

an LVDT electronic displacement sensor, and dividing by the gage

length. The stiffness of the load train, relative to the sample,

yields an accuracy in these measurements of ± 0.04 percent. Load

was applied by a 2-kN servohydraulic testing and measured by a compatible load cell in series with the pull rod. Load was measured

to within ± 0.2 percent.

The specimens were strained at a constant rate of 0.002

cm/cm/min using the testing system's interval function generator

and a linear ramp wave form. All of the specimens were strained

between 0.025 and 0.040 at 4 K. Some experiments were conducted

by sequentially turning the magnetic field on and off during the

test. In such tests, the specimen was unloaded, but kept at 4 K,

78
Fig. 2. Schematic of applied tensile load versus specimen strain

with magnetic field applied "on", then "off".

prior to field change. A schematic of a typical applied loadspecimen-strain curve during these tests is shown in Fig. 2.

To determine the percent (bcc, a') martensite in the specimens

at room temperature following deformation, a bar-magnet with a torsion spring was-y^ed. This apparatus has been previously described

and calibrated. ' The wire specimens had a small diameter com- pared to the bar magnet. To increase the sensitivity of the measurements, the entire 27.5-cm length of the deformed specimens was cut

into smaller lengths, about 2-cm long. Therefore, the magnetic

measurements represent an entire specimen average, with sections

included over which the field strength varied from 0 to 7 T.

Amounts of martensite transformation were estimated. from the magnetic

readings, using earlier established conversions.

EXPERIMENTAL RESULTS

Experimental data for alloy 310 are tabulated in Table I, and

the experimental data for alloy 304 are summarized in Table II.

In these tables, ANN refers to annealed, and CD refers to cold

79
Table

I.

AISI

310

Stainless

Steel

Wire*

0>

C

4J o

*H •H c

4-) c CO

0> 00 o O

O o o o O

O

Q

O o O 4J Co

03 r—

z:

w

Cfl 4-J 3

CO »4H o oo o 3 C

+-

C *H vO

m f-H

C 4-> •H -3 1 1 1 1 CNI NT CN| •H <D 4-1 rH 1 1 1 1 • • • * •K * * *

<0 Cfl c a f-H _H i-H c o *h

*-» o u >< CO Cfl *HQ

trength f

fset), 3 o in

m o O o o O o O o o CO O Cl- o o 00 ON CN| CO o CO i-H O o

z: nt CO CNI CNI <—* CNI CN o f-H o o O

Yield (0.2%

f-H

rH

COco -

*H 4-1 U -H CO sO >3-

m o s£> CNI CNI 00 00

m

m

a Cl- 00 co f-H f-H 00

m

m p»

m 0 -H z •—

rH co sD CO CNI CNI NT NT NT

O-O

»h Ph

Cfl Cfl - 3

00 rH 03 vO NT CN| 00 n£> v£> CNI rH '-O P- o C 3 Ph 00 P- VO vO O O o p** P- 00 00 00 3 X) o CN| CNI CNI O O

>-

X c p o

<0 *H a 4-1

Q 2 Q

Q

Q 2 2 2 2 2

•H *H CJ ++ <^5 2 CJ CJ

a 2 2 2 2 2 U 03 < < < < < < 0) C a o

CO

V

u

•H -

4-1 *3

QJ rH C 0) 00 *H

CO Ch

H

O o OO

O

OO P*» p^ p^ Pz: *

0)M3

4-J 4-)

Cfl 03 in inm in CNJ CNI CN| CN| CN| CN| CN| CN| 3 V- ON O' ON o> H <D CNI CN| CNI CN| <r NT NT 'T NT NT NT NT ae3H

U

<U

QJB

T3I

v£> O* CBm

p-MO

Uh

u:ON

COm

T3

Q> rH

03(1) cc<

++ c

•HBmuo

4-t

ON

co

co

4J

03 T3

<D rH

03

<UCc<

U7>

*3" oo

Mho

Cfl c

•H

03U

Cfl

'O

0)>

80 t

Total

strain

(elastic/plastic).
drawn (as received) . Small amounts of martensite were detected in

all alloy 304 specimens strained at 76 and 4 K, but not in the

alloy 304 specimens strained at room temperature. No martensite

was detected in alloy 310 specimens. In Table II, the martensite

(a’) values are reported as percent martensite divided by the

specimen elongation; it has been previously found that the amount

of a ' is approximately linear with specimen strain.

Constant-field measurements

There was no detectable effect of a 7-T magnetic field on the

yield strength at 4 K of either the 304 or 310 alloys. However,

the data scatter is large, and the application of the magnetic field

would not necessarily produce changes of flow strength in excess

of the scatter of yield strength values. This confirms the findings

of Bolshutkin, Desnenko, and Ilichev'' who also could not detect

any distinction between zero field and 3.4-T field tensile properties at 4 K for a series of Fe-18Cr-8 to 25Ni alloys.

The magnetic field did not significantly affect the initiation

of discontinuous yielding in the cold-drawn 310 alloy. No discontinuous yielding was observed in the annealed 310 alloy to strains of

0.04. In contrast to the 310 alloy behavior, serrated yielding was observed at lower strains for annealed alloy 304, compared with a cold-drawn condition. Data related to the initiation of serrated

yielding for alloy 304 (in terms of strains), contained in Table II, are not sufficiently -reproducible to permit assessment of the role

of applied magnetic fields.

Field switching measurements

Experiments were performed in switching the magnetic field

from zero to 7 T and from 7 T to zero after the specimens had been

partially deformed at 4 K and then unloaded and maintained at 4 K

(see Fig. 2). On subsequent reloading after field switching no change of flow strength was observed in either alloy 304 or 310.

That is, after application of field, there was no accompanying detectable increase or decrease of flow strength at 4 K. The measurement equipment was capable of detecting changes of about 0.2 percent

in flow strength.

During the sequential on-off magnetic field tensile tests,

the Young's modulus was measured while the specimen was unloaded

and loaded. While absolute values varied for alloy 304 by about

± 13 percent for five separate specimens, the variation of each

measurement on the same specimen is estimated to be ± 1 percent.

The initial data were obtained after the first unloading, usually

at plastic strains of 0.001 to 0.003. Subsequent measurements were

obtained at field switchings, usually in’ strain increments of 0.002

to 0.003. The elastic modulus was found to decrease linearly with

81
•K

<U •Hs

<y

01 4-> oo

cn

Cfi 0) r—

c

•H

to

4-1 cn<rom

c/3 M

<c

0)

r-H

cd H

<D

Co

•H •H W u c CO nt <—

--4 CN| •—I

m CO 'O 'D 00

03 00 o o

u>

C o o O CN CN f—l r—4

u o

00 r-H x w

5^

CO

4-J 3

03 ^ o oo O 3 C

4-

C -H o sD o O'

m CO v£>

O CO 1—4 «—

C w *h -a 1 1 1 vO m 1—

CO

m v£> 00 o>m r-- 00

•H 0) 4-> rH 1 1 1 • • • 03 tfl C a) f—l -H 1—4 O O O O o o O o C O *H

4->

O

u

C/3 CO *HQ• U /-s 00 4J C <13 OJ CO U *4-4 U *4-1 OJ CO 00 o o CO

m 00

O O' NT CO CO

C/3

O P-4 CNI CO 00 nt r- CN CO 00 CN •—4 CN 00 £ CO CNI CO <T s£> m m NT m \D m NT TJ 6^

r-H CN

03 . •H O >, w

rH

03Co * •H 4-» 4J *H 03 v£> 00 00 NT NT in

m O' 00 O'*

4- E P.

O O' 00 CO CO CNI CN NT NT

O CN v£> NT vO

O -H a. g o

>-•

X

pi

CO CO • 3

00 rH 03 vO 00 nt \D CO CNI CN CN O' r- NT o O' C 3 Pm O' O o •—4 O' r-H r—

•—i o —

i—

CO CN 3 HD O c—

CN 1—4 CM CNI <—CN CN CM CN CN CN CN CN

O

O

>< Z C e o

03 •H E 4J

a z z a

QQ 2 z Z z Z z XX *H •H cj 2 cj o CJ Z z z z z z Z U -a < < 3 < 3 <d <2 <

a> c

a. o

(73 CJ

u

•H -

4_» -a

03 f—

. C 03 00 *H H o C" O o C" i-** o o o r» r*» r^. o

nj U.

—

X O O •> 03U3

4-1 4-» CO 03

m

m vD cn CNI CN CN CN CN CN CM CN CN CN

03

U

u O' ON r- • H 0) CN Csi >3- NT NT NT NT NT NT NT NT

QuE

03 H

v0)

O)6

03 •H

*0I <5- \DO*

'Ocoue

*Heo

4- 44

T3

r-H

OJ

82

All

values

of

Young's

modulus

were

taken

at

0.5-

1.0%

strain

by

unloading

and

loading.
increasing plastic deformation in the range between 0.001 to 0.020

strain for both 304 and 310 alloys.

Because the absolute values of the Young's modulus had considerable scatter, the data were normalized to 0.010 strain to values

of 210 GPa for alloy 304 and 206 GPa for alloy 310 at 4.2 K. ' The adjusted Young's modulus data are plotted as a function of

strain in Fig. 3 for both zero and 7-T fields. There is no distinction between zero and 7-T field measurements of Young's modulus.

Both tend to scatter approximately equally above and below the

linear trend line for alloy 304. There is a clear linear dependence of Young's modulus on strain within the strain interval to

0.020. Finally, the Young's modulus for alloy 310 depends slightly

less on plastic deformation than for alloy 304.

Fig. 3. Dependence of Young's modulus (E) , measured at 4.2 K, on specimen strain at 4.2 K for alloys 304 and 310 with zero

and 7-T magnetic fields. Note that the absolute values of

E are "normalized" to 210 GPa for alloy 304 and 205 GPa

for alloy 310 at a strain of 0.010.

83
DISCUSSION

No variation in Young's modulus (E) with the application of a 7-T transverse magnetic field could be measured. Our Young's

modulus measurement sensitivity between tests on a single specimen

is estimated to be one percent. Considering that alloy 304 Young's

modulus is about 210 GPa, and the measurement is made at flow

stresses of about^600 MPa or less, the attendant total specimen

strain is 3 x 10 '. Changes of E of the order of 2.1 GPa

(1 percent) would be reflected in changes of strain of 3 x 10

It follows that the transverse magnetostriction of both alloys must

be less than about ± 3 x 10 for our measurement system not to

detect changes of E.

Similarly, it is unlikely that the magnetostriction of the

body-centered cubic (a') martensite phase contributes to distinctions between zero and 7-T field alloy 304 tensile measurements.

Only about two percent martensite was detected. Transverse magnetostriction strain contributions of the a' phase would need Jjo be of

the order of 3 x 10 per 2 percent martensite = 1.5 x 10 to

change E by one percent. Changes of body-centered cubic and

face-centered cubic transverse magnetostriction coefficients at

roo^ temperature are reported by Bozorth^ not to exceed 0.02 x

10 throughout the Fe-Ni composition range.

The flow strength could possibly be influenced by magnetic

field enhancement of a' transformation that, in turn, affects specimen flow strength. The enhancement of a' formation in Fe-Ni and

Fe-Ni-Mn alloy systems has been discussed earlier and, therefore,

it is likely that some minimum field strength, field orientation,

and/or pulse frequency would produce more martensite transformation.

Increased strain-induced a' transformation is thought to play a significant role in mod|.|^c^ion of the low temperature stressstrain characteristics. ’ However, neither the martensite

enhancement nor the modification of stress-strain behavior have

been detected under the conditions of this experiment.

Finally, the dependences of the Young's modulus on plastic

deformation at 4.2 K of both alloys are of interest and have not

been previously reported. The linear reduction may either be

caused by defect production (at 4.2 K neither point nor line defects

can annihilate), by the effect of applied stress or defect production on antiferromagnetic domains or spin-glass clusters, or by

dislocation unpinning. It would be of fundamental interest to

accurately assess, at 4 K, the roles of applied stress and plastic

deformation on elastic constants of austenitic steels.

84
REFERENCES

1. Korenko, M. K. and Cohen, M., "Martensitic transformation in

high magnetic fields," in Proc. Conf. on Martensitic Transformations , ICOMAT 1979, Massachusetts Institute of Technology,

Cambridge, MA (1981), pp. 388-393.

2. Fokima, Y. A., Smirnov, L. V. , and Sadovskiy, V. D., "Effect

of a pulsating magnetic field on the position of the martensitic transformation range in steel," Fiz. Met. Metalloved.,

Vol . 19, No. 4 (1965), pp. 592-595.

3. Fokima, Y. A., Smirnov, L. V. , and Sadovskiy, V. D., "Destabili

zation of austenite by a strong pulsating magnetic field," Fiz.

Met. Metalloved., Vol. 19, No. 5 (1965), pp. 722-725.

4. Fakidov, I. G., Voronchikhin , L. D., Zavadskiy, E. A., and

Barkhanov, A. M. , "The variation in the magnetic intensity of

austenitic steel exposed to a strong pulsating magnetic field,

Fiz. Met. Metalloved., Vol. 19, No. 6 (1965), pp. 852-857.

5. Vornochikhin, L. V. and Fakidov, I. G., "Determining the latent

heat of the martensitic transformation induced by a magnetic

field in steel," Fiz. Met. Metalloved., Vol. 21, No. 3 (1966),

pp. 436-441.

6. Malinen, P. A., Sadovskiy, V. D., and Sorokin, I. P., "Effect

of magnetic field on the martensitic transformation in alloys

of the N G, type," Fiz. Met. Metalloved., Vol. 24, No. 2 (1967), pp. 305-309.

7. Voronchikhin, L. D., Romashev, L. N., and Fakidov, I. G.

,

"Influence of the initial state of austenitic steel on the

martensitic transformation in a strong magnetic field," Fiz.

Met. Metalloved., Vol. 26, No. 5 (1968), pp. 948-951.

8. Malinen, P. A. and Sadovskiy, V. D., "Effect of magnetic field

on the e-a transformation in Fe-Mn alloys," Fiz. Met.

Metalloved., Vol. 28, No. 6 (1969), pp. 1012-1017.

9. Bolshutkin, D. N., Desnenko, V. A., and Ilichev, V. Ya., "Low

temperature plastic deformation of Fe-Cr-Ni steels in magnetic

fields to 2.71 x 10 AM-1," Cryogenics, Vol. 19 (1979),

pp. 231-234.

10. Gotishcheve, V. I., Glinik, R. A., Petrovskii, M. L., and

Khazov, V. N., "Effect of magnetic field on the plastic deforma

tion of aluminum at 4.2 K," JETP Letters, Vol. 30, No. 2 (1979)

pp. 92-96.

11. Lebedev, V. P., Khotkevich, V. I., Krylovskiy, V. S., and

Philips, A. H. , "Deformation of nickel in a steady magnetic

field," Fiz. Met. Metalloved., Vol. 48, No. 5 (1979), pp. 1 1 08—

1111 . 12. Reed, R. P. and Tobler, R. L., "Deformation of metastable

austenitic steels at low temperatures," R. P. Reed and A. F. Clark, eds., in Advances in Cryogenic Engineering-Materials , Vol. 28, Plenum Press, New York (1982), pp. 49-56.

85
13. Collings, E. W. and King, H. W. , "Austenitic stainless steels

as magnetic transition metal alloys, E. W. Collings and H. W.

King, eds., in The Metal Science of Stainless Steels , Metallurgical Society (AIME) , New York (1978), pp. 1-21.

14. Ekin, J. W., "Fatigue and stress effects in NbTi and Nb^Sn

multifilamentary superconductors," in Advances in Cryogenic

Engineering-Materials , Vol. 24, Plenum Press, New York (1978),

pp. 306-316.

15. Reed, R. P. and Guntner, C. J., "Sttess-induced martensitic

transformations in 18Cr-8Ni Steel," Trans. AIME, Vol. 230

(1964), pp. 1713-1720.

16. Ledbetter, H. M. , "Stainless-steel elastic constants at low

temperatures," J. Appl. Phys., Vol. 52 (1981), pp. 1587-1589.

17. Bozorth, R. M. , Ferromagnetism , Van Nostrand Co., New York

(1951), p. 669.

86
FACTORS INFLUENCING THE LOW-TEMPERATURE DEPENDENCE OF YIELDING

IN AISI 316 STAINLESS STEELS*

R. L. Tobler, D. H. Beekman, and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Tensile tests at temperatures between 323 and 4 K were performed on one heat of AISI 316 austenitic stainless steel having the composition Fe-17.34Cr12. 1 7 N

i - 1.55Mn-2. 16Mo-0.051C. The temperature dependences of the yield and

flow strengths at plastic stain increments from 0.2 to 3.65 % are analyzed.

At the yield strain (0.2 %), no body-centered cubic (bcc) martensite phase

transformation is detected. At higher strains (“3.2 ± 0.6 %), bcc martensite

forms from the parent austenite phase at test temperatures below 175 K, but

there are no discontinuities in the temperature dependence of flow strength.

A review of data available for three heats of AISI 316 at temperatures between

973 and 4 K reveals that deviations from thermally-activated plastic flow

theory occur at temperatures below 175 K, apparently depending on heat-to-heat

compositional variations. Grain size and magnetic transition effects on the

yield strength are discussed.

To be published in Austenitic Steels at Low Temperatures, Eds. R. P. Reed

and T. Horiuchi, Plenum Press, New York (1983).

87
'
INTRODUCTION

The mechanical behavior of the AISI 300 series austenitic

stainless steels over a broad range of temperatures is a subject of

interest to designers of superconducting and nuclear energy devices.

In this paper, factors influencing the initial yielding and stressstrain behavior of AISI 316 stainless steel are described. AISI 316

stainless steel is a metastable Fel6-18Cr, 10-14Ni, lMo steel that

may undergo a partial austenite-to-martensite phase transformation

during deformation at low temperatures, the probability of which

increases with increasing strain, decreasing temperature, and

decreasing alloy content.

When transformation occurs in Fe-Cr-Ni austenitic stainless

steels, two martensitic phases may be formed: the hexagonal closepacked (hep) e martensite and the body-centered cubic (bcc) a'

martensite. The e phase usually forms first, then a* appears [1].

The e phase is difficult to detect in small quantities, whereas the

a' phase is ferromagnetic and easy to detect using a magnetometer [2], For AISI 316, it is reported that no a' forms spontaneously on cooling to 4 K, but up to 50% a* forms on straining to

failure at 4 K [3]

.

The published data [4-9] for AISI 316 indicate substantial

heat-to-heat variations in yield strength at cryogenic temperatures,

presumably due to the broad range of chemical composition allowed in

the material specifications. The explanation of such variation is

complicated because there has been no systematic study and investigators have not always reported the complete details of material

89
characterization and test procedure. In particular, the effects of

grain size and martensitic phase transformation have never been

thoroughly evaluated at cryogenic temperatures.

Soviet authors [10,11] report the existence of three anomalies

in the temperature dependence of the yield strength for Fe-18Cr-xNi

austenitic stainless steels containing 8, 10, or 20% Ni. These

anomalies amount to sharp decreases of the yield strength on the

order of 100 MPa, occurring over narrow temperature intervals. The

first anomaly occurs between 20 and 4 K and is unexplained. The

second anomaly occurs at temperatures between 60 and 35 K and is

attributed to the Ndel transition (magnetic ordering) . The third

anomaly occurs in metastable alloys (lower Ni) at temperatures in

the martensitic phase transformation range. One aim of this paper

is to determine whether such anomalies are evidenced in AISI 316.

Of particular interest is the initiation of martensitic phase

transformation, because there is apparently conflicting evidence on its role in affecting the low-temperature yielding and flow behavior

of metastable stainless steels. Suzuki et al. [12], measuring the

temperature dependence of the 0.2% offset yield strength in an Fe-18Cr-9Ni stainless steel, report a significant decrease in flow

strength when the test temperature is lowered from about 260 K.

They argue that e martensite forms at lower temperatures, reducing

the flow strength. Following e martensite formation, a' martensite

forms at slip-plane intersections and the a' lathes along <110> act

as windows for dislocation pileup release; the result is an "easy-glide" (work-softening) region of the stress-strain curve. This "easy-glide" region had previously been reported by Reed and

Guntner [13], who argued that it was caused by strain-induced e martensite, since their x-ray and transmission electron microscopy

data for an Fe-18Cr-8Ni alloy indicated an increasing concentration

of e during easy glide. In contrast, Olson and Azrin [14], studying

an Fe-9Cr-8Ni alloy, found an excellent correlation between flow

strength reduction at lower temperatures and the initiation of a*

martensite. In their experiments, a' martensite was induced by

elastic stresses at low temperatures with attendant reduction of

yield strength. The low-temperature martensite transformation was suggested to be stress assisted and to originate from existing

nucleation sites. No "easy-glide" region of the stress-strain curve was observed.

To study the above effects, continuous flow-strength measurements as a function of temperature are described here, as an extension of previous work [4]. AISI 316 stainless steel was selected

because it exhibits strain and temperature ranges where no martensite forms, while at lower temperatures, body-centered cubic

martensite forms at larger strains.

90
MATERIALS

The AISI 316 stainless steel tested in this study was obtained

in the form of a 38-mm-thick plate. Tensile specimens were

machined in the longitudinal orientation in the form of flat

coupons having a 50.8-mm gage length and a 2.54-mm x 4.77-mm-gage

cross section. The machined specimens were annealed at 1338 K for

0.5 h, producing a Rockwell B (R^) hardness of 71 and an average

grain diameter of 55 ym (ASTM grain size number 5.5). The chemical

compositions of this steel (denoted heat 2) and other AISI 316

steels referred in this paper are listed in Table 1.

PROCEDURE

Tensile tests were performed on (1) as-machined and annealed

or (2) strained and recrystallized material. The strained and

recrystallized material was used to determine grain-size effects.

All tensile tests were performed at a crosshead rate of 8.5 ym/s

using a screwdriven machine and cryostat [1,16]. Test temperatures of 4 and 76 K were achieved with baths of liquid helium and

liquid nitrogen. Other temperatures were achieved by controlled

thermal conduction using liquid nitrogen reservoirs and heating

wires attached to the specimen grips and load train. Current to

the heating wires was varied until thermal equilibrium was reached

at the specified test temperature, which was accurate to ± 1.5 K.

Load-versus-def lection curves were recorded with a commercial

load cell and clip-gage extensometer . Clip-gage extensometers

permit measurement and magnification of the specimen strain within

the (2.5-cm) gage length. Estimated flow-strength measurement

inaccuracies are ± 1%, but specimen-to-specimen variations

Table 1. Chemical composition (weight percent) for AISI 316

steel heats referred to in this study.

Heat No. & Reference

Fe Cr Ni Mn C p s Si Cu Mo Co N

1 [A] balance 17.25 13. A8 1.86 0.057 0.024 0.019 0.58 2.34 — 0.03

2 this study II 17.3A 12.17 1.55 0.051 0.026 0.014 0.68 0.34 2.16 0.19

3 [5] II 16.72 11.27 1.52 0.05 — — 0.32 2.58

A [15] It 17.46 12.93 1.77 0.062 0.030 0.020 0.56 2.20

91
including both material and measurement variations) may reach

± 5%. As an indication of data scatter, seven tests were performed on the AISI 316 steel at 76 K, resulting in an average

yield strength (a ) of 448 MPa and an overall spread of ± 7 MPa.

After plastic strains of 0.002 to 0.038 (mm/mm), the specimens were unloaded, warmed to 295 K, and examined for the presence of

ferromagnetic a' martensite using a torsion-bar magnetometer [2],

X-ray diffractometer measurements were used to try to detect small

concentrations of e martensite. No e was detected and our estimated sensitivity is 2%.

Grain-size variations were obtained by recrystallization

treatments. Tensile specimens were plastically strained 20 to 30%

at 295 or 76 K and then annealed for 0.25 to 0.5 h in an argon

atmosphere at 1233 to 1573 K. These treatments produced grain

sizes ranging from 22 to 140 ym. After the recrystallization

treatments, the specimen surfaces were cleaned by electropolishing. The recrystallization treatments increased the scatter of

yield-strength measurements at 76 K to about ± 15 MPa. Texture

effects were neglected.

A few experiments were performed on the annealed AISI 316

stainless steel to evaluate the strain rate effect. Increasing

the crosshead rate by a hundredfold from 0.85 ym/s to 85 ym/s increased the flow strength at strains lower than 0.01 by increments

of 6% at 295 K and 3% at 76 K. The effect at 4 K was negligible.

Therefore, strain rate does not play a significant part in explaining the variability of results noted in the text.

The M temperature and strain-hardening coefficients were measured after loading a series of specimens to 0.027 or 0.037

plastic strain at low temperatures. The percent a' was measured

at 295 K using a magnetometer; then the M^ temperature was obtained

by extrapolating a plot of a* per unit strain versus temperature.

The strain-hardening coefficients (n) at each test temperature were measured as the slope of the log a-versus-log e plots. Such

plots are nearly linear, but only for small strain increments.

The values of M^ and n apply only for the strain levels studied in

this paper.

RESULTS

The tensile yield strength (o ) and initial stress-strain behavior of the Fe-17Cr-12Ni-2Mo all^y were measured as a function

of temperature. Grain size, martensitic transformation, and

strain-hardening effects are described in this section.

92
Table 2. Yield strength measurements at 0.002 plastic strain

obtained for AISI 316 (heat 2) of this study.

Test Temperature, T,

K

Yield Strength, o , MPa* y

323 185

295 204

295 198

295 198

240 244

215 281

194 298

192 306

190 307

164 349

154 368

143 380

125 399

100 428

76 456

76 451

76 445

76 450

76 441

76 447

76 448

60 456

39 487

23 528

4 537

4 527

*To convert from MPa to ksi, divide by 6.895.

Temperature

Twenty-six a measurements were obtained for annealed AISI 316

stainless steel (Keat 2) at seventeen temperatures between 323 and

4 K (Table 2) . Magnetometer inspections of the specimens indicated

no a' martensitic transformation for this heat after 0.002 plastic

strain, regardless of test temperature. The o data between 295

and 4 K are plotted and compared with previously published data

[4-9] in Fig. 1. This figure illustrates significant heat-to-heat

variations of a , with considerably greater differences at cryogenic temperatures than at room temperature.

93
Fig. 1. Yield strength of AISI 316 stainless steels at low temperatures, illustrating heat-to-heat variability. Heat

1 is from Tobler et al. [4]; heat 3 is from Sanderson and

Llewellyn [5]. Other data: o - Read and Reed [6]; O - Baughman [?]; v - Tobler et al. [8]; - Hoke et al . [ 9 ]

To evaluate the data it is first necessary to characterize

the temperature dependence. Several references cited in Fig. 1 report data at only a few test temperatures, whereas the functional dependence of o on temperature can be properly identified

only by extensive measurements over a broad temperature range.

Therefore, relevant data for three heats of AISI 316 are replotted

in Fig. 2, which shows that a increases progressively as temperature is lowered. Despite min^r variations of grain size, strain

rate, and composition, the a values for the three heats exhibit

parallel trends between 922 ^nd 175 K, where the test-temperature

intervals from different studies overlap. Below 175 K, significant

variations of strength and cross-over effects begin to emerge. At

still lower temperatures, the variations become more obvious

until, at 76 K, o correlates with composition: Fe-17.25Cr-13.47Ni

(heat 1) exhibits^the highest strength, Fe-17.34Cr-12.17Ni (heat 2)

exhibits intermediate strength, and Fe-16.72Cr-ll.27Ni (heat 3)

exhibits the lowest strength. The evidence that this is a true

effect of alloy composition is considered in a later section.

94
TEMPERATURE, K

Fig. 2. Yield strength data for AISI 316 stainless steels in the

elevated-to-cryogenic temperature range, illustrating

overall temperature dependence. The symbols a and a

are defined in the Discussion. Heat 1 - Tobler et al.

[4]

; heat 2 - this study; heat 3 - Sanderson and Llewellyn

[5]

.

Strength variations may arise between specimens from a single

heat of steel owing to inhomogeneity or thermomechanical processing

effects. As shown in Fig. 1, Read and Reed's data [6] for a 316

stainless steel at 295, 76, and 4 K are 5 to 17% higher than the

results of Tobler et al. [4], who tested the same heat. In the

former study, the specimens were tested in the as-machined condition, whereas the specimens in the second study were re-annealed

(1289 K, 1 h) after machining. This step reduced the hardness at

295 K from R^79 to R^75 and lowered a by increasing the grain

size slightly and removing any residual cold work from the fabrication or machining processes. This effect was also observed by

Sikka [17].

Grain size

Figure 3 illustrates the effect of grain size on a at 295,

76, and 4 K. The data at each temperature fit an inverse square

root dependence on the average grain diameter, d, in reasonable

accord with the Hall-Petch equation:

o=o. + kd (1)

y i

95
Here a. refers to the flow strength in the absence of grain

boundaries (fhe "friction stress") and k is the slope of the

a -versus-d 2 plot and indicates the grain boundary dependence of

tKe yield strength. The values of o. and k are given in the

figure for each test temperature (29^, 76, and 4 K)

Although the grain size effect is significant at low temperatures, it alone is not sufficient to account for the wide variation in properties for different heats. For example, heat 1 (d = 67 pm) has a higher yield strength than heat 2 (d = 55 pm) at 4 K,

despite its larger grain size. This supports the discussion to

follow: that a significant portion of the strength variations at

4 K is due to alloy composition.

Martensitic transformation

Specimens (heat 2) deformed at low temperatures were measured

using a magnetometer, but no a' martensite was detected after

0.002 plastic strain. To induce measurable transformation, a series of specimens were loaded to higher strains at 76 K.

Figure 4 presents the measured data of the strain dependence of

a’. A plastic strain of about 0.022 is necessary for a* to form

Fig. 3. Effect of grain size on the yield strength at three test

temperatures for the Fe-17Cr-12Ni-2Mo alloy.

96
at 76 K. The dependence of a' on plastic deformation is linear,

at least to 0.05 strain. At larger strains, the dependence

apparently increases (more a' forms per unit strain) since fractured specimens show between 44 and 52% a' at uniform strains of

0.4 to 0.5.

The relationships between flow stress, plastic deformation,

and temperature are presented in Figs. 5 and 6, and Fig. 7 illustrates the experimental method of determination. From

Fig. 7, the quantity of a' martensite formed after 0.037 plastic

strain is negligible for test temperatures between 323 and 175 K.

Below 175 K, a' forms in increasing amounts, and the percent a*

martensite per unit strain is linearly dependent on test temperature. The curves of Figs. 5 and 6 represent specimens in various

degrees of transformation, containing from 0 to 10% a* martensite

per unit strain, but the stress-strain curves are regular and free

from discontinuities. In other words, we see that the martensitic

transformation in this steel apparently occurs gradually with no anomalous decrease of strength. These findings contrast with the

behavior of Fe-18Cr-8Ni and Fe-18Cr-10Ni stainless steels that

exhibit: (1) irregular yield and flow strength temperature

dependences owing to the initiation of e and a' phase transformations [10], and (2) suppression of a

1 martensite at 4 and 20 K,

compared with that at 76 K [13].

Fig. 4. Dependence of bcc (a’) martensite on tensile plastic

strain. Martensite was detected by using a bar magnet

and torsion balance beam.

97
Fig. 5. Temperature dependence of the yield and flow stresses

for AISI 316 stainless steels, showing no evidence of

anomalies despite martensitic phase transformation

occurring at higher strain increments and lower test

temperatures

Strain-hardening behavior

Two parameters of strain hardening, n and Ao, are plotted

versus temperature in Fig. 8. A simple and convenient measure of

strain hardening is Ao, taken as the difference in flow strength at

plastic strains of 0.0275 and 0.002 (Ao = o^

o275

_°0 002^' ^he

Ao remains constant between 295 and 175 K, then increases below

175 K. From Figs. 7 and 8 we see corresponding trends for Ao and

the a’ transformation: both increase linearly with temperature

below 175 K. The strain-hardening coefficient, n, exhibits a minimum when plotted versus temperature, and it also begins to

increase at temperatures near 175 K. We conclude that the onset of

a' transformation in AISI 316 causes the increased strain

hardening, and that the degree of strain hardening is proportional

to the amount of a’ martensite formed at these low strains.

98
Fig. 6. Stress-strain curves for AISI 316 stainless steels at' various test temperatures, showing no evidence of anomalies despite martensitic phase transformations occurring

at higher strain increments and low temperatures.

DISCUSSION

Thermally activated flow

The behavior of a for AISI 316 stainless steels is consistent

with equations based o^ thermally activated plastic flow. According to Seeger [18], o may be separated into two components:

ay °A

+ ( 2 )

Here, the athermal component, a , is a long-range stress-field

contribution in which the stresses are large compared with thermal

fluctuations and where a^/G (G = shear modulus) is independent of

temperature. Other dislocations on parallel slip planes may

contribute to a . The thermal component, a , is a short-range

stress-field contribution in which the obstacles to dislocation

movement may be overcome with the aid of thermal fluctuations.

99
Fig. 7. Amount of bcc martensite per unit strain in AISI 316

specimens deformed to plastic strains of 0.0275 or 0.0365, indicating clear evidence of martensitic phase

transformation at cryogenic temperatures. Extrapolated

temperature of 0% bcc martensite is about 175 K.

Contributions to the short-range stress fields come from forest

dislocations, solute atoms, small jogs in dislocations, and the

Peierls-Nabarro stress [19].

In Fig. 9 the o data for the three AISI 316 steels have been

normalized to the sh^ar modulus and plotted versus temperature,

which is normalized to the melting point, = 1658 K. (Shear

modulus data at cryogenic temperatures were taken from Ledbetter

[15] and at elevated temperatures from Moore [20]). These data

indicate that o^ becomes active at temperatures lower than about

1/3 T^. The divergences of the thermal components at low temperatures are observable in Fig. 2.

The following expression derives from thermodynamics and

thermally activated plastic flow theory [21]:

°T

= (°

0 “ °A

) exP(“CT ) (3)

100
Here is the extrapolated value of o at zero temperature and C

is a materials constant related to activation energy. A complete

expression of o is then:

y

°y

- °A

+

<°0 • °A) exP(-CT > (4)

In a study of alternative equations and curve-fitting techniques, Tobler showed that Eq. (4) best characterized the a - versus-T data for AISI 316 [4]. In this equation, the values

require no correction for the shear modulus, but are estimated

from a at elevated temperatures (Fig. 2). The estimates for

heats X, 2, and 3 are 100, 95, and 113 MPa, respectively.

Data that agree with Eq. (4) exhibit a linear trend when

plotted using semilogarithmic coordinates, assuming a single

rate-controlling thermally activated plastic flow mechanism is

operative over the entire temperature interval. Or, if different

processes are rate controlling for specific temperature intervals,

the trend will consist of several linear segments.

c

UJz

o

0. X

UJo

z

z

UJo

QC<

X

Z

<

<r

fr- CC

Fig. 8. Temperature dependence of strain-hardening parameters.

101
TEST TEMPERATURE, K T_ MELTING TEMPERATURE, k’ Tm

Fig. 9. Normalized yield strength versus normalized temperature

for three AISI 316 stainless steel heats, showing variations in thermal and athermal stress components.

As illustrated in Fig. 10, Eq. (4) does adequately characterize the data at elevated and cryogenic temperatures. The data for

all three heats can be fit using bilinear trends, where a slope

transition occurs at about 60, 150, or 175 K, depending on the

heat

The main trend of data for heat 1 fits the expression:

o y100 + 700 exp (-6 . 2 x 10

3 T) (5)

Deviation from this equation begins at about 60 K, reflecting a change in the exponential.

One consequence of the slope transition is that is lower

(17 and 33% for heats 1 and 2) than would otherwise be the case.

102
Fig. 10. Temperature dependence of the thermal-stress component,

illustrating applicability of thermally activated plastic flow theory at temperatures between 900 and 175 K.

This has practical implications for alloy design. For example, in

applications requiring maximum strength at extreme cryogenic

temperatures, it would be desirable to suppress or eliminate the

inflection point in In o for these steels. Our data indicate

that this can be done by raising the Ni content to the maximum

allowable limit of the material specification.

What can be said about the cause of the slope transitions in

Fig. 10? It was suggested previously that dislocation tunneling,

magnetic ordering, and adiabatic heating processes may affect o at temperatures approaching absolute zero [4], But, since thes^

processes are not operative above 60 K, they cannot account for

the deviations observed in heats 2 and 3 at 150 and 175 K.

A more plausible explanation is that the deviations represent

a change in mechanism of plastic deformation, which relates to

alloy composition. As noted in Fig. 10, the deviations in

for heats 1, 2, and 3 occur at about 60, 150, and 175 K,

respectively. This sequence follows the order of decreasing alloy

103
content (Table 1), decreasing stacking fault energy (Table 3), and

decreasing crystal structure stability, as indicated by the M and

temperatures (Table 3). Therefore, it is possible that a

S secdnd thermally activated deformation mechanism having a lower

activation energy becomes rate controlling at low temperatures.

If so, a change of slopes as in Fig. 10 would be expected. This

could be the case in the present study if faulting of the austenite

occurred at low temperature, or if e phase formed in quantities

too low to be detected. The possible martensite roles are discussed in the next section.

Role of martensitic transformation

Martensitic transformations may have the following effects on

the low-temperature deformation of AISI 316 alloys: (1) they may

be responsible for the slope transitions in plots of In versus

Table 3. Calculated stacking fault

T^ temperatures for AISI

energies

316 steel

(SFE) and

heats of

M , M , and

this study.

Heat No. & Reference SFE,

3

mJ/M2 v

b

K

M

C V

K

T

d

K

1 [4] 32.3 -226 190 24.7

2 this study 29.3 -126 232 25.5

3 [5] 27.8 - 47 246 7.4

4 [19] NA NA NA 26.4

a. Room temperature values, calculated from equation (1) of

Rhodes and Thompson [22]. (The SFE is expected to decrease at

lower temperatures)

b. M is the temperature at which martensite starts forming

spontaneously on cooling, as calculated from the equation of

Larbalestier and King [23].

c. is the temperature above which no deformation induced martensite is possible, as calculated from the equation of

Williams et al. [24].

d. T^ is the temperature of the onset of antiferromagnetic

ordering, as calculated from the equation of Warnes and

King [25].

104
temperature below about 175 K, and (2) they may influence the

stress-strain behavior at sufficient strain (>0.02, Fig. 4) and

low temperature (>175 K, Fig. 7).

To provide a relative indication of the austenite stability,

the M

g and M^ temperatures (Table 3) were calculated on the basis

of the Cr, Ni, Mn, Si, C, and Mo contents. The calculated M

g values are negative, indicating complete stability on cooling to

absolute zero, which is consistent with the fact that no a' martensite was detected after cooling to 4 K. The calculated M and

temperatures are only approximate and relative guides to the

behavior of the AISI 316 steels at cryogenic temperatures.

Comparison of the measured and calculated M values for heat 2 shows considerable disparity, the calculated value being 57

degrees too high. Absolute agreement is not achieved because the

predictive equations were developed on lower alloy steels, in

compression tests, for specific levels of martensite, and without

consideration of trace element effects. Moreover the M temperature is rather loosely defined, because high strains will produce

some transformation at higher temperatures (see the 195-K datum.

Fig. 7).

First, it is unlikely that the a’ transformation can account

for the observed slope transitions (Fig. 10), since no a’ martensite was detected at the yield strength (0.002 plastic strain).

From Fig. 4 it is apparent that no a* is expected to form at 76 K

until a plastic strain of about 0.02 is reached. However, it is

possible that martensitic transformation involving small amounts

of e phase or austenite faulting prior to a* transformation may be

responsible. From Table 3, the calculated stacking-fault energies

decrease, but remain positive, with increased deviation of a^, as illustrated in Fig. 10. This possibility could not be confirmed,

however, since our metallographic and x-ray inspection methods

lacked the required sensitivity to detect less than 2% of e.

Second, martensite transformations- usually influence deformation characteristics. In the commercial Fe-18Cr-8Ni austenitic

stainless steel (AISI 304), the e and a' martensitic products

apparently contribute to the anomalous temperature dependence of

a and to the "easy-glide" portion of the stress-strain curves at

temperatures near 77 K [12,13,26], However, Figs. 4 through 7 suggest that in the alloy, of this study (Fe-17Cr-12Ni-2Mo) , the

temperature dependence of the flow strength (including o ) monotonically increases with strain, and therefore, the a' martensitic

product does not cause anomalous behavior.

With respect to the role of martensitic products in affecting

the temperature dependence of the flow strength of Fe-Cr-Ni alloys

105
at low temperatures, there are two schools of thought: Suzuki

et al. [12] argue that the e phase causes a reduction of a at low

temperatures since in their study a’ martensite was detected only

at stresses considerably greater than a . They have argued from

metallographic evidence that e martensite forms at stress levels

near the measured a . Additionally, during the easy-glide portion

of the stress-strait curve for their Fe-18Cr-8Ni alloy at low

temperatures, they detect a’ martensite and find that it typically

forms at slip-band intersections. The reduced flow strength in

the easy-glide region is attributed to the strain-induced a' transformation that forms at slip-band intersections and acts as "windows" for dislocation pileups. Conversely, Olson and Azrin

[14] studied a less stable Fe-9Cr-8Ni alloy and obtained excellent

correlation between the temperature dependence of a and the flow

strength at which a' martensite was first detected.

From our study, it is apparent that the onset of straininduced a' martensite in austenitic Fe-Cr-Ni alloys does not

necessarily lead to a reduction of the flow strength at low

temperatures. Rather, the formation of a' in our Fe-17Cr-12Ni

steel served simply to cause higher rates of strain hardening.

Therefore, in the Fe-Cr-Ni alloy system, great significance

attaches to alloy composition, and especially to Ni contents in

the range 8 to 14%. This finding raises the question of the

nature of the strain-induced martensitic product or products that

form(s) during plastic deformation. Conflicting parametric

studies and lack of conclusive microscopy at this time preclude a final assessment of the role of e and a' martensites in affecting

low-temperature flow strength.

The relationship between the "normalized" percent bcc martensite (percent magnetic martensite divided by plastic strain) and

temperature (Fig. 7) deserves additional experimental clarification. It is implied that the temperature of about 175 K represents the onset of strain-induced martensite, and at higher

strains and temperatures little magnetic product phase is detected.

This behavior was substantiated by a few tests to fracture at

195 and 295 K (Fig. 7). However, the nature of the family of curves that could be generated by taking data at larger strains

remains undefined. Similar data for AISI 304 steel indicate the

existence of a constant temperature of a' initiation [27], but a gradual deviation from linearity at larger strains.

Role of grain size

Data on the grain-size effect on the yield strength of

austenitic stainless steels are rare. The room temperature data

of this study indicate a very slight effect, as evidenced by the

106
TM"2 low k value of 191 MPa'yrn^ . This value is in reasonable agreement with the 221 MPa*ym 2 value obtained by Norstrom [28] for a 316L stainless steel. Low k values are typical for face-centered

cubic alloys at room temperature. At cryogenic temperatures,

however, the grain-size effect increases significantly. As shown

in Fig. 3, the values of a. and k at 4 K are greater by factors of

2 and 5, respectively, than the values at 295 K.

In Fig. 11, the temperature dependences of k and a. from

Eq. (1) are plotted, using the 4- to 295-K data from this study

combined with the 295- to 875-K data of Norstrom [28]. The k is

strongly temperature dependent at low temperatures, but insensitive to temperature above room temperature. The stress, a.,

increases markedly with decreasing temperature for temperatures

lower than about 500 K. These data clearly imply that grain size

has a significant effect on low-temperature flow strength. From

Fig. 11. Dependence of the tensile yield strength, o ; the HallPetch parameter, k; and the friction stress^ o., for

AISI 316 steel. High-temperature a. data from Norstrom

[28] are adjusted to 0.04 weight percent nitrogen.

107
Fig. 11, the quantity (o - a.) represents the grain-boundary

contribution to flow strength and about 30% of a at 4 K is

contributed by grain-boundary strengthening.

Role of magnetic transition

Similar to other Fe-Cr-Ni austenitic stainless steels,

AISI 316 undergoes a magnetic ordering phenomenon characterized by

a change from the paramagnetic to the antiferromagnetic state at

low temperatures. Accordingly, a peak occurs in the temperature

dependence of magnetic susceptibility at T , the Ndel transition

temperature, which corresponds to the onset of antiferromagnetic

ordering.

Since a decrease of elastic constants occurs at [29,30],

it was suggested that the Ndel transition also affects dislocation

motion and hence tensile properties [10,11]. On the basis of

Soviet experimental data [10,11] for related steels, a sharp

decrease of a of about 100 MPa would be expected for our Fe17Cr-12Ni all^y at T . N

Fig. 12. Temperature dependence of yield strength and Young's

modulus for AISI 316 stainless steels, showing no drop

in yield strength at the Ndel transition temperature.

108
As shown in Fig. 12, data for an AISI 316 (Fe-17.5Cr-12.9Ni)

stainless steel do reveal a measurable drop in Young’s modulus

(E) , identifying the Ndel transition between 40 and 20 K. But the

data of Fig. 12 indicate no decrease of a corresponding to the

magnetic transition at T^. Nor is there ^ny evidence of an anomalous decrease of a at temperatures between 20 and 4 K. So

this study does not support the conclusions of other investigators, and no explanation for the apparent conflict can be offered

at this point.

SUMMARY AND CONCLUSIONS

Tensile tests at temperatures between 323 and 4 K were

performed on metastable austenitic AISI 316 stainless steel

(Fe-17Cr-12Ni-2Mo) . The yield and flow strengths at plastic

strains up to 0.037 are analyzed and compared with existing data,

noting metallurgical effects. The grain-size effect on yield

strength (a ) is weak at room temperature but substantial at

cryogenic temperatures. No martensite is detected after tensile

loading at low temperatures to 0.002 plastic strain. With further

straining (^0.03), body-centered cubic (a*) martensite forms in

the parent austenite phase at test temperatures below 175 K. The

initiation of a’ transformation in this steel in tension is

associated with increased strain hardening below 175 K, and not

with a decrease of flow strength. Significant features of the

temperature dependence of flow strength are: 1) a continuous rise

of a with decreasing temperature, in accord with thermally

activated plastic deformation, 2) an inflection in the thermal

component of a at about 150 K, suggesting a change in deformation

mechanism, and^3) absence of anomalies due to martensitic transformations, or to magnetic transition at Ndel temperature.

Comparison of present results with data for other heats indicates

a significant variation due to variations in chemical compositions

within the material specifications.

ACKNOWLEDGMENTS

We thank D. Beard, Office of Fusion Energy, DoE for his

support and continued encouragement. D. Burkhalter of NBS contributed significantly to this study by conducting tensile tests.

Dr. E. L. Brown of NBS, who contributed valuable metallographic

assistance and reviews of the manuscript, is also gratefully

acknowledged.

109
REFERENCES

1. Read, D. T. , Reed, R. P. , and Schramm, R. E. , Low temperature

deformation of Fe-18Cr-8Ni steels, in: Materials Studies for

Magnetic Fusion Energy Applications at Low Temperatures-II , F. R. Fickett and R. P. Reed, eds., NBSIR 79-1609, National

Bureau of Standards, Boulder, Colorado (June 1979),

pp. 151-172.

2. Reed, R. P. and Mikesell, R. P., The stability of austenitic

stainless steels at low temperatures as determined by

magnetic measurements, in: Advances in Cryogenic

Engineering , Vol. 4, Plenum Press, New York (1960),

pp. 84-100.

3. Larbalestier , D. C. and King, H. W. , Austenitic stainless

steels at cryogenic temperatures I, structural stability and

magnetic properties. Cryogenics 13:160-168 (1973).

4. Tobler, R. L., Reed, R. P. , and Burkhalter, D. S., Temperature

dependence of yielding in austenitic stainless steels, in:

Advances in Cryogenic Engineering , Vol. 26, Plenum Press,

New York (1980), pp. 107-113.

5. Sanderson, G. P. and Llewellyn, D. T., Mechanical properties of

standard austenitic stainless steels in the temperature range

-196 to +800C, J. Iron Steel Inst. 207:1129-1140 (1969).

6. Read, D. T. and Reed, R. P., Fracture and strength properties

of selected austenitic stainless steels at cryogenic temperatures, Cryogenics 21:415-417 (1981).

7. Baughman, R. A., Gas atmosphere effects on materials. Progress

Report #2, AF 33(616), General Electric Co., Schenectady,

New York (November 1958).

8. Tobler, R. L., Mikesell, R. P., Durcholz, R. L., Fowlkes,

C. W. , and Reed, R. P. , Fatigue and Fracture Toughness

Testing at Cryogenic Temperatures , NBSIR 74-359, National

Bureau of Standards, Boulder, Colorado (March 1974), pp. 182—

308.

9. Hoke, J. H., Mabus, P. G., and Goller, G. N., Mechanical

properties of stainless steels at subzero temperatures. Met

Prog. 55:643-648 (1949).

10. Ilyichev, V. Ya., Medvedev, Ya. M. , Shapovaliev, I. A., and

Klimenko, I. N. , Low temperature anomaly of the temperature

dependence at the flow stresses in iron-chroraium-nickel

alloys, Phys. Met. Metallogr. 44(1) : 173-176 (1978).

11. Verkin, B. I., Ilyichev, V. Ya., and Klimenko, I. M. , The

low-temperature change of the magnetic structure and plastic

properties of Fe-Cr-Ni alloys, in: Advances in Cryogenic

Engineering, Vol. 26, Plenum Press, New York (1980), pp. 120-

125.

12. Suzuki, T., Kojima, H., Suzuki, K. , Hashimoto, T. , and

Ichichara, M. , An experimental study of the martensite

nucleation and growth in 18/8 stainless steel, Acta Metall. , 23:1151-1162 (1972).

110
13 . Reed, R. P. and Guntner, C. J., Stress-induced martensitic

transformations in 18Cr-8Ni steel, Trans. AIME , 230:1713-1720

(1969)

.

14. Olson, G. B. and Azrin, M. , Transformation behavior of TRIP

steels, Metall. Trans. A , 9A:713-721 (1978).

15. Ledbetter, H. M. , Stainless-steel elastic constants at low

temperatures, J. Appl. Phys . 52 (3) : 1587-1589 (1981).

16. Reed, R. P., A cryostat for tensile tests in the temperature

range 300 to 4 K, in: Advances in Cryogenic Engineering , Vol. 7, Plenum Press, New York (1961), pp. 448-454.

17. Sikka, V. K. , Tensile and creep properties of type 316

stainless steels, paper CONF 78 1219-3 (N80-11212), presented

at the UK/UKAEA Specialists Meeting on Mechanical Properties,

Oak Ridge, Tennessee (Dec. 4, 1978).

18. Seeger, A., The generation of lattice defects by moving

dislocations, and its application to the temperature dependence of the flow-stress of FCC crystals, Philos. Mag.,

46:1194-1217 (1955).

19. Conrad, H., The cryogenic properties of metals, in: HighStrength Materials , V. F. Zackay, ed., Wiley, New York

(1964), pp. 436-509.

20. Moore, T. D. , Structural Alloys Handbook , Vol. 2, Battelle’s

Columbus Laboratories, Columbus, Ohio (1980).

21. Yaroshevich, V. D. and Ryvkina, D. G., Thermal-activation

nature of plastic deformation in metals, Sov. Phys. -Solid

State (Engl. Transl.) 12 (2) : 363-370 (1970).

22. Rhodes, C. G. and Thompson, A. W. , The composition dependence

of stacking fault energy in austenitic stainless steels,

Metall. Trans. 8A: 1901-1906 (1977).

23. Larbalestier , D. C. and King, H. W. , Prediction of the low

temperature stability of type 304 stainless steel from a room temperature deformation test, in: Proceedings of the Fourth

International Cryogenic Engineering Conference , K. Mendelssohn, ed., IPC Science and Technology Press,

Guildford, Surrey, England (1975), pp. 338-340.

24. Williams, I., Williams, R. G., and Capellaro, R. C., Stability

of austenitic stainless steels between 4 K and 373 K, in:

Proceedings of the Sixth International Cryogenic Engineering

Conference , ed., K. Mendelssohn, IPC Science and Technology

Press, Guildford, Surrey, England (1976), pp. 337-341.

25. Warnes, L. A. and King, H. W. , The low temperature magnetic

properties of austenitic Fe-Cr-Ni alloys 2, the prediction of

Ndel temperatures and maximum susceptibilities. Cryogenics

16(11) : 659-667 (1976).

26. Guntner, C. J. and Reed, R. P., The effect of experimental

variables including the martensitic transformation on the low

temperature mechanical properties of austenitic stainless

steels, Trans. Am. Soc. Met. , 55:339-419 (1962).

27. Starr, C. D. , Notes on the plastic critical temperature in

strain induced martensite reaction, Trans. AIME 197:654 (1953).

Ill
28 . Norstrom, L. A., The influence of nitrogen and grain size on yield strength in Type 316L austenitic stainless steel. Met.

Sci. 11:208-212 (1977).

29. Collings, E. W. and Ledbetter, H. M. , Sound velocity anomalies

near the spin glass transition in an austenitic stainless steel

alloy, Phys. Lett. 72A( 1 ) : 53-56 (1979).

30. Ledbetter, H. M. and Collings, E. W. , Low-temperature magnetic

and elastic-constant anomalies in three manganese stainless

steels, in: The Metal Science of Stainless Steels , E. W. Collings and H. W. King, eds.. Met Soc. AIME, New York

(1978), pp. 22-38.

112
TEMPERATURE DEPENDENCE OF FLOW STRENGTH OF

SELECTED AUSTENITIC STAINLESS STEELS+*

R. P. Reed, R. L. Tobler, and J. W. Elmer

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

It is important to characterize and to understand the temperature dependence of the tensile flow strength of austenitic steels for efficient design

and material selection. Recent studies have suggested the possibility of

three anomalies in the temperature dependence of the flow strength of Fe-Cr-Ni

and Fe-Cr-Ni-Mn austenitic stainless steels. Reduction of flow strength at

decreasing temperature may be associated with the onset of the austenite to

martensite transformation (^200 K), the magnetic transition at the Ndel

temperature (^50 K) and low temperature dislocation dynamics ( <20 K). The

tensile flow strength was investigated in the temperature range 4 to 295 K for

two annealed polycrystal 1 ine alloys: Fe-20Cr-16Ni-6Mn-0.2N steel (produced in

the USSR and independently tested by two research laboratories) and Fe-18CrlONi-O.lN steel (AISI 304LN). The former alloy is stable with respect to

strain-induced martensitic transformations, the latter metastable.

+ Work supported by Office of Fusion Energy, Department of Energy

* To be published in Austenitic Steels at Low Temperatures, Eds. R. P. reed

and T. Horiuchi, Plenum Press, New York (1983).

113
•' •
INTRODUCTION

Austenitic stainless steels are used extensively as the

structural material in cryogenic systems, and knowledge of the

dependence of their strength on temperature is sometimes

essential for proper design. This study addresses the

temperature (T) dependence of the tensile flow strength in the

range 4 K to room temperature.

Recent studies [1-3] of Fe-Cr-Ni and Fe-Cr-Ni-Mn austenitic

stainless steels have reported three anomalies in the temperature

dependence of the flow strength. These anomalies consist of

sharp decreases of strength on the order of 100 MPa occurring

over narrow temperature intervals between room temperature and

4 K. One type occurs only in metastable steels and is associated

with the onset of austenite-to-martensite transformation, usually

at about 200 K. For both stable and metastable steel compositions, two additional anomalies are reported at temperatures

between 60 and 35 K and between 20 and 4 K. The flow strength

decrease in the temperature interval 60 to 35 K has been attributed to the paramagnetic-to-antif erromagnetic transition at

the Ndel temperature, where the elastic constants are known to

decrease. It has been suggested that flow strength reductions

below 20 K result from possible dislocation tunneling mechanisms [2], Studies in our laboratory [4-6] have identified the

decrease in flow strength associated with austenite instability

*This research was supported by the Office of Fusion Energy, U.S.

Department of Energy.

115
with respect to martensitic transformations. This paper reports

on the temperature dependence of the flow strength of two austenitic steels over the entire temperature range between 4 and

300 K. Data are presented for an Fe-20Cr-l 6Ni-6Mn-0 . 2N steel

(produced in the USSR and independently tested by two research

laboratories) and for an Fe-18Cr-10Ni-l . 5Mn-0 . IN steel (AISI

304LN) . The former alloy is expected to be stable with respect

to martensitic transformations, the latter metastable.

MATERIALS

A 25-mm plate of nitrogen-strengthened austenitic stainless

steel was obtained from the Paton Institute of Electrowelding,

Kiev, USSR. The material was tested in the as-received

condition. The characteristics of the alloy are listed here.

Composition (wt.%): Fe-20.3Cr-16.4Ni-6.lMn-0.19N-0.06C-0.2lMo0. 26Si-0.025P-0.026S-0.09Cu. Hardness: Rockwell B88, at 295 K

after machining tensile specimens. Microstructure : Equiaxed

austenitic grains, free of carbonitride precipitation when viewed

at 400X. Grain size: Average grain diameter of 0.039 mm (ASTM

No . 6.4).

The AISI 304LN material was received from a U.S. supplier.

It was one of nine alloys produced from three 136-kg vacuuminduction-melted heats, split with respect to carbon level, and

teemed into 76-cm x 200-cm x 360-cm hot-topped cast-iron ingot

molds. The ingots were then reheated and soaked at 1561 K,

hot-rolled to 25-mm-thick plates and air cooled. The plate was given a final anneal at 1332 K for one hour and water quenched.

The characteristics of the alloy are listed here.

Composition (wt.%): Fe-18.0Cr-9.9lNi-l.50Mn-0.016P-0.017S0. 5 7 Si-0. 205Mo-0. 197Cu-0. 037C-0. 12N. Hardness: Rockwell B79 at

295 K after machining tensile specimens. Microstructure:

Equiaxed austenitic grains. Grain Size: Average grain diameter

of 0.105 mm.

PROCEDURE

In a previous study, unpolished, longitudinally oriented

specimens were used to measure tensile yield (o ) and ultimate

strengths, as well as ductility, at selected temperatures [ 5 ]

In this study, transverse-oriented specimens were used for yield

strength measurements at variable temperatures. Both round

(6.25 mm diameter) and flat (6.25 mm x 2.50 mm) specimens with

25-mm gage length were used. They were electropolished in

perchloric acid solution (70% methanol, 25% glycerol, 5%

perchloric acid) to remove scratches and cold-worked material

from machining operations. A total of 0.1 mm (4 mils) was

116
removed from the diameter, which is about four times scratch

depth.

Tensile tests us^d a screw-driven machine at a crosshead

velocity of 8.5 x 10 mm/s, with apparatus previously described

[4,5,7], Variable temperatures were achieved using controlled

thermal conduction from cryogenic fluid reservoirs attached to

the load train and electric heaters on the specimen grips. The

automatic regulators maintained temperatures to ± 0.2 K, with an estimated absolute uncertainty of ± 1 K [5]. A strain-gage

extensometer was clipped on the specimen at a gage length of

25 mm.

Specimen-to-specimen scatter of the temperature dependence

of the flow strength was eliminated by performing temperaturechange tests. To obtain measurements at three to five temperatures, a single specimen was loaded to a strain within the range

0.0005 to 0.002 at T^, then unloaded, cooled to T^, and reloaded

until plastic flow resumed. This procedure is described schematically in Figure 1. Cooling to the lower temperature usually

took 10 to 20 minutes. Additional loadings at lower temperatures

were also performed, and the flow strength obtained by extrapolating to the 0.002 strain.

€

Figure 1. Schematic of test procedures for multiple temperature

measurements of flow strength using one specimen.

117
RESULTS

Magnetic measurements [4] at 295 K after low temperature

(76, 4 K) deformation confirmed that the Fe-20Cr-16Ni-6Mn-0. 2N

alloy does not exhibit any a' (body-centered cubic) martensitic

transformation and that the Fe-18Cr-10Ni-l . 5Mn-0. IN alloy is

metastable and partially transformed to a* martensite. This was expected since calculations listed in Jable 1 of T [8] (the

temperature at which a* martensite is expected tom

^>egin to

transform on cooling), and T [9] (the highest temperature , of

deformation- induced a' martensite) are both quite low for the

Fe-20Cr-16Ni-6Mn-0 . 2N alloy and T for the Fe-18Cr-10Ni-l . 5Mn0.1N alloy is sufficiently high. Stacking fault energy (SFE)

estimations [10], listed also in Table 1, indicate that the

Fe-20Cr-16Ni-6Mn-0. 2N alloy is relatively more stable with

respect to the face-centered cubic to hexagonal close-packed (e)

strain-induced martensitic transformation.

Alloy Fe-20Cr-16Ni-6Mn-0. 2N

The yield strength results are presented in Figure 2, which

compares data from this study with those of our previous

study [5] and with the results of Yushchenko et al. [1]. Between

300 and 4 K, the results show approximate conformation to an exponential function given:

ay = 1160 exp (-4.16 x 10~ 3 T) (1)

Most of the data fit this equation to within about 6 percent,

although two points deviate by larger amounts (-10 percent at

295 K, +8 percent at 60 K)

Figure 3 shows only the data at test temperatures between 4

and 100 K. For these results, combined with our previous data,

the data spread for five tests at 4 K is ± 55 MPa, owing

primarily to specimen-to-specimen variability. The previously

tested specimens [5] were longitudinal and unpolished, but the

results are indistinguishable from the electropolished transverse

specimens

Table 1. Austenite Stability and Calculations of Stacking Fault

Energy.

00 i 1 T

[9]

md

(K)

SFE^

10 ^ Alloy ms

(K) (10

-5

J/m2

)

Fe-20Cr-16Ni-6Mn-0.2N -715 18 79

Fe-18Cr-10Ni-1.5Mn-0. IN - 84 285 28

118
Figure 2. Yield strength versus temperature for Fe-20Cr-16Ni6Mn-0.2N steel, summarizing data from this study and

other studies [1-3,5].

Temperature, K

Figure 3. Low temperature yield strength data for alloy Fe-20Cr16Ni-6Mn-0 . 2N. Data points connected by lines represent yield strengths obtained from single specimens.

Each specimen was initially strained at highest temperature, then cooled and restrained.

119
The 7-percent drop of a at T , as reported by Yushchenko

et al. [1], falls well withi^i specimen-to-specimen variability

determined in our study.

To eliminate the obscuring effect of specimen-to-specimen

variability, single-specimen temperature-change tests were

performed. In Figure 3, the multiple data points from single

specimens are linked with straight lines. These trend lines have

nearly equivalent slopes, in general do not cross over one another, and include the temperature range (50-65 K) in which

strength decreases at lower temperatures have been re- ported [ 1-3 ]

Alloy Fe-18Cr-10Ni-l . 5Mn-0 . IN

Tensile flow strength data at strain offsets of 0.0005 and

0.002 (yield strength) are plotted as a function of temperature

in Figure 4. The difference between the flow strengths at 0.002

Figure 4. Flow strength at strains of 0.0005 and 0.002 as a function of temperature for Fe-18Cr-10Ni-l . 5Mn-0. IN

alloy. Solid line represents average temperature

dependence from individual specimen measurements,

reported in Figure 4.

120
and 0.0005 strains increases from about 0.038 GPa at 295 K to

about 0.070 GPa at 4 K.

Between 300 and 200 K, the yield strength best conforms to:

ay

= 900 exp (-4.22 x 10~ 3 T) (2)

The limited data restrict the assignment of data inaccuracies to

this function. Between 4 and 200 K, the dependence of a on T

decreases to: ^

ay

= 730 exp (-3.05 x 10

_3

T) (3)

The data fit this equation within ± 5 percent.

Temperature-change experiments were also performed on the

Fe-18Cr-10Ni alloy. In Figure 5, the measurements from seven specimens are plotted. For all measurements, the temperature was

lowered after reaching strains between 0.0005 and 0.002. The

temperature dependence found from single-specimen tests is

plotted as a smooth curve in Figure 5. All specimens exhibited

normally rising flow strength with a decrease in temperature,

with the possible exception of the specimen originally strained

at 295 K. The change of strength achieved from the changing of

temperature matches that obtained from individual tests at the

two temperatures.

Figure 5. Flow strength of Fe-18Cr-10Ni-l .5Mn-0. IN alloy at

strain of 0.002 (yield strength) measured from speci- mens sequentially cooled and strained.

121
Temperature, K

Figure 6. Comparison of data from Verkin et al. [3] and this

study for similar Fe-18Cr-10Ni alloys.

Our data and data for a similar Fe-18Cr-10Ni-lMn alloy [3] are presented for comparison in Figure 6. Three major ranges of

strength decrease with decreasing temperature are reported by

Verkin et al. [3]. None of the anomalous strength decreases are evident from our measurements on an alloy that is similar in base

composition and in total interstitial content. The specimen-tospecimen data scatter for our Fe-18Cr-10Ni-l . 5Mn-0 . IN alloy is

much less than the anomalous decreases reported for the Fe-18CrlONi-lMn alloy [3],

DISCUSSION

From Figure 4 the dependence of the flow strength on temperature of the Fe-18Cr-10Ni-l . 5Mn-0 . IN alloy may be characterized

by a change of deformation mechanism below about 200 K, producing

two linear regions (4 to 200 K, 200 to 300 K) in the log a-T

graph. The temperature change experiments (Figures 3 and 5) do

not conflict with this characterization. The austenite of the

Fe-20Cr- 16Ni-6Mn-0 . 2N alloy is relatively very stable and has

higher stacking fault energy and lower T (below 0 K) than the

Fe-18Cr-10Ni-l .5Mn-0. IN alloy, owing to higher alloying concentrations (Table 1). It is likely that the strain-induced

martensitic transformation (to hexagonal close-packed e or a')

leads to the change in the temperature dependence of the flow

122
strength below 200 K in the less stable alloy Fe-18Cr-10Ni-l . 5Mn0.1N. The a* martensite was detected in Fe-18Cr-10Ni-l . 5Mn-0 . IN

specimens deformed at 76 and 4 K, and the estimated T for this

alloy is calculated to be positive (Table 1). Conversely, no a*

martensite transformation was detected in the Fe-20Cr-16Ni-6Mn0.2N alloy and only one linear region of log a-T is observed over

the entire 4 to 295 K temperature range. However, since e and a*

transformation are closely associated in austenitic stainless

steels, it is not clear from these data which martensitic product

may be associated with flow strength changes in the Fe-18Cr-10Ni1.5Mn-0.1N alloy.

The single-specimen temperature-change experiment data have

been presented in Figures 3 and 5. Clearly there is no distinction at low temperatures (<200 K) in the temperature dependence

of the flow strength determined from multiple or single specimen

measurements. This indicates that the deformation mechanisms for

plastic strain and the defect structures generated during plastic

strain are essentially independent of temperature for these two

alloys. That is, the same flow strength at a given temperature

and strain can be achieved either by straining at the prescribed

temperature or by straining to a lower strain at higher temperatures (but <200 K) , cooling to the prescribed temperature and

straining. This result is a bit unpredictable, since the low

temperature deformation mechanisms of the less stable austenitic

steel (Fe-18Cr-10Ni-l . 5Mn-0 . IN) include martensitic transformation. In this discussion, the strain range is limited to <0.002;

at larger strains with more complex martensitic transformation,

the flow strength may not be (strain) path independent.

Finally, this study has failed to confirm the flow strength

decreases associated with the Ndel transition temperature. The

flow strength (a ) is proportional to the shear modulus (G) and

the dislocation density (p) through the relationship

a

f a Gb (4)

where a, a material constant, usually has a value between 0.2 and

0.5 (e.g., Conrad [11]). The elastic constants for Fe-20Cr-16Ni6Mn-0.2N steel [12] show a regular, nearly linear increase as temperature is reduced from 300 K, until a plateau is approached

below about 76 K. In the 60 to 48 K range, the E and G moduli

decrease by about \ percent. This decrease marks the Ndel

temperature, T , at which there is a change in magnetic ordering

from paramagnetic to antiferromagnetic conditions. The midpoint

of the transition is about 54 K. The ^-percent decrease of G

measured at T for the Fe-20Cr-16Ni-6Mn-0. 2N alloy would not

•N alone be sufficient to account for the 7-percent decrease of a , as reported (Figure 2) by Yushchenko et al. [1] and, for an ^ Fe-18Cr-10Ni-lMn alloy, the 17-percent decrease reported by

123
Verkin et al. [3]. Indeed, the ^-percent effect is indistinguishable from experimental scatter in o measurements. The

single-specimen temperature change test Conducted below 60 K

shows an increase of o with a magnitude equivalent to that

expected from the exponential function of Eqs. (1) and (3).

CONCLUSIONS

Two austenitic stainless steel alloys have been measured in

tension between 4 and 300 K. One alloy contains 20Cr, 16Ni, 6Mn,

and 0.2N; the other contains 18Cr, lONi, 1.5Mn, and 0.1N and is an A1SI 304LN. The results are summarized below:

1. The temperature dependence of the flow strength of the

two alloys is best characterized by a linear dependence

of log a versus T. The more stable Fe-20Cr-16Ni-0 . 2N

alloy exhibits a linear slope of log a versus T over

the entire temperature test range (4 to 295 K) . The

less stable Fe-18Cr-10Ni-l . 5Mn-0 . IN alloy exhibits two

linear ranges on a log a versus T plot, one from 4 to

about 200 K, the other from about 200 to 295 K.

2. Temperature change experiments at low strains (^0.002)

indicate reversibility of flow strength, at temperatures

below about 100 K. This implies that the deformation

mechanisms do not change under these conditions for

these alloys.

3. The reported anomalous low temperature dependences of

the flow strength associated with the Ndel transition

temperature and very low temperatures (<20 K) for

austenitic steels were not confirmed by this study.

REFERENCES

1. Yushchenko, K. A., Verkin, R. I., Ilichev, V. Ya. , and

Klimenko, I. N., "The influence of magnetic structure on temperature dependence of the yield strength of 20Cr-16Ni6Kn steel at low temperatures," in Advances in Cryogenic

Engineering-Materials, Vol. 28, Plenum Press, New York,

pp. 67-72 (1982).

2. Ilichev, V. Ya., Medvedev, Ya. M. , Shapovalov, I. A., and

Klimenko, I. N. "Low temperature anomaly of the temperature

dependence of the flow stresses in iron-chromium-nickel

alloys," Fiz. Met. Metalloved. Vol. 44, No. 2, pp. 173-176

(1980)

124
3. Verkin, B. I., Ilichev, V. Ya. , and Klimenko, I. N. , "The

low temperature change of the magnetic structure and plastic

properties of Fe-Cr-Ni alloys," in Advances in Cryogenic

Engineering-Materials , Vol. 26, Plenum Press, New York,

pp. 120-125 (1980).

4. Reed, R. P. and Simon, N. J., editors. Material Studies for

Magnetic Fusion Energy Applications at Low Temperatures - IV , National Bureau of Standards, Boulder, CO, NBSIR 81-1645

(1981)

.

5. Reed, R. P. , Tobler, R. L. , Elmer, J. W. , McHenry, H. I.,

and Yushchenko, K. A., "Strength and toughness of USSR

Fe-20Cr-16Ni-6Mn-0 . 2N weldments at cryogenic temperatures,"

in Proceedings of the Eighth International Cryogenic Engineering Conference , IPC Science and Technology Press,

Guildford, Surrey, England, pp. 797-801 (1980).

6. Tobler, R. L., Beekman, D. H., and Reed, R. P., "Lowtemperature dependence of yielding in AISI 316 stainless

steel" in Austenitic Steels at Low Temperatures , R. P. Reed

and T. Horiuchi, editors, Plenum Press, New York (1982).

7. Read, D. T. and Tobler, R. L., "Mechanical property measurements at low temperatures," in Advances in Cryogenic

Engineering-Materials , Vol. 28, R. P. Reed and A. F. Clark,

editors, Plenum- Press , New York (1982).

8. Eichelman, G. H. and Hull, F. C., "The effect of composition

on the temperature of spontaneous transformation of austenite to martensite in 18-8-type stainless steel," Trans. Am.

Soc. Met., Vol. 45, pp. 77-104 (1953).

9. Williams, I., Williams, R. G., and Capellero, R. C., "Stability of austenitic stainless steels between 4 K and

373 K," in Proceedings of the Sixth International Cryogenic

Engineering Conference , IPC Science and Technology Press,

Guildford, Surrey, England, pp. 337-341 (1976).

10. Schramm, R. E. and Reed, R. P., "Stacking fault energies of seven commercial austenitic stainless steels," Metall.

Trans. Vol. 6A, pp. 1345-1351 (1975).

11. Conrad, H. , "The cryogenic properties of metals," in HighStrength Materials , V. F. Zackay, Editor, John Wiley & Sons,

New York, pp. 436-509 (1965).

125
12. Ledbetter, H. M. and Austin, M. W. , "Low-temperature elastic

properties of an Fe-20-16Ni-6Mn alloy," in Material Studies

for Magnetic Fusion Energy Applications at Low Temperatures - V , R. P. Reed and N. J. Simon, editors, National

Bureau of Standards, Boulder, CO, NBSIR 82-1667 (1982).

126
PROPERTIES OF AUSTENITIC STAINLESS STEEL AT CRYOGENIC TEMPERATURES+*

H. I. McHenry

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

The low temperature behavior of austenitic stainless steels is reviewed

with emphasis on three general classes: 1) the AISI 300-series in the an- nealed condition; 2) the nitrogen-strengthened grades; and 3) high-strength

cold-rolled sheet of the AISI 300-series. Mechanical and physical properties

of selected alloys in these classes are presented. Welding, the properties of

weldments, and other fabrication considerations are discussed. Recent results

on the influence of ferrite content, nitrogen, and sensitization on castings

of composition similar to the AISI 300 series are also reviewed.

+ Work supported by Office of Fusion Energy, Department of Energy

* To be published in Austenitic Steels at Low Temperatures, Eds. R. P. Reed

and T. Horiuchi, Plenum Press, New York (1983).

127

INTRODUCTION

The austenitic stainless steels are Fe-Cr alloys with sufficient

nickel and manganese (and sometimes nitrogen) to stabilize austenite,

a face-centered cubic (f.c.c.) phase. Chromium provides corrosion

resistance. The most typical austenitic stainless steels are variations of the 18Cr-9Ni alloy. Although none of the commercial alloys

were specifically developed for low-temperature service, several of

them have been widely used in a variety of cryogenic applications.

Their popularity stems from their retention of excellent mechanical

properties, particularly toughness, at low temperatures coupled with

ready availability, ease of fabrication, and good service experience.

Their disadvantages are they are more expensive and have lower yield

strength than ferritic Fe-Ni steels and aluminum alloys, and their

machinability is poorer than that of aluminum alloys.

The austenitic stainless steels have a wide range of mechanical

properties, depending on their chemical composition and thermalmechanical processing. Three general classes of these alloys are useful for cryogenic applications:

a. The AISI 300-series stainless steels in the annealed condition, particularly types 304 and 304L, are the most commonly used

alloys. Consequently, they have the most service experience and

coverage in design codes. These grades have moderate strength and

excellent toughness, and they are readily available in a variety of

product forms. Casting alloys equivalent to the commonly used

300-series grades are used to manufacture complex shapes, such as valve bodies.

129
b. The nitrogen-strengthened grades are remarkably stronger at

cryogenic temperatures and retain most of the advantages of the 300-

series grades, except for availability, service experience, and code

coverage

c . The cold-rolled sheets of the AISI 300-series grades have

been used as high-strength sheet materials in specialized applications,

such as liquid-fueled missile cases.

In addition to the three main classes, there are precipitationhardened austenitic grades, for example ASTM A 453, which offer promise

as high-strength alloys for heavier-gauge applications, but their usage

to date has been limited to nuts and bolts. Precipitation-hardened

alloys must be selected carefully because these alloys are generally

martensitic or semiaustenitic and have low toughness at cryogenic

temperatures

ANNEALED AISI 300-SERIES GRADES

The designations and compositions of the most commonly used grades

of the AISI 300 series are summarized in Table 1. These grades have

moderate strength, excellent toughness, and good fabrication characteristics, and they are readily available in a variety of product forms.

basic

The alloys listed in Table 1 can be compared with AISI-304, the

19Cr-9Ni stainless steel. Type 316 contains 2 to 3% Mo to improve

Table 1. Compositions of austenitic stainless steels.

AISI

Type No.

Composition (%)

Cr Ni C, max. Other

301 16-18 6-8 0.15

302 17-19 8-10 0.15

304 18-20 8-12 0.08

304L 18-20 8-12 0.03

305 17-19 10.5-13 0.12

309 22-24 12-15 0.20

310 24-26 19-22 0.25 1.5 Si, max. 310S 24-26 19-22 0.08 1.5 Si, max.

316 16-18 10-14 0.08 2-3 Mo

316L 16-18 10-14 0.03 2-3 Mo

321 17-19 9-12 0.08 (5 x %C) Ti, min.

347 17-19 9-13 0.08 (10 x ZC) Nb+Ta, min.

130
resistance to pitting in chloride solutions, a property that led to

its selection for cargo piping systems aboard liquef ied-natural-gas

(LNG) ships. Types 304L and 316L are low-carbon modifications of 304

and 316, respectively; low carbon is desirable to avoid sensitization,

the grain-boundary precipitation of chromium carbides. Types 321

(alloyed with titanium) and 347 (alloyed with niobium and tantalum)

are also sensitization-resistant grades. Type 310S is a 25Cr-20Ni

alloy that is a low-carbon version (0.08^ C max.) of AISI Type 310

(0.25% C max.). The high-alloy content provides austenite stability,

and thus 310S is useful where dimensional stability and nonmagnetic

behavior are essential.

The properties of the 300-series austenitic stainless steels have

been compiled in cryogenic handbooks, 1 ’ 2 in a review article by Brickner

and Defilippi, 3 and in suppliers publications. 5 In this section,

the general trends in strength, toughness, and fatigue resistance are summarized as a function of temperature. Use the data presented here

and the data in the handbooks with care because there may be significant

heat-to-heat variations in the mechanical properties of the 300-series

stainless steels at low temperatures, particularly in the yield strength

and toughness. The principal causes of these variations are interstitial content (carbon and nitrogen) and mechanical deformation history, which controls the resulting grain size. The nitrogen content,

even within the specification limit of 0.08 wt.%, is particularly

significant because higher nitrogen causes substantial increases in

yield strength.

Tensile Properties

The ultimate tensile and yield strengths of AISI grades 304, 310,

and 316 are compared in Fig. 1 for temperatures ranging from 4 to

300 K. 6 Notice that type 304 has the lowest yield strength and the

highest ultimate strength at all temperatures. The increased ultimate

tensile strength of type 304 is due to the martensitic transformation,

which occurs more readily in 304 than in ,316 or 310. Since the transformation is strain induced, the yield strength is not influenced by

martensite, but the ultimate strength is increased.

The ultimate tensile and yield strengths of AISI grade 304L are slightly lower than those of grade 304. The loss of strength associated

with reduced carbon content is considered to be characteristic of all

the stainless steels (e.g., 304L vs. 304, 316L vs. 316, and 310S vs.

310) and is attributed to the strengthening effect of carbon.

The ultimate tensile and yield strengths of AISI grades 304, 321,

and 347 are essentially the same, indicating that the small additions

of titanium in type 321 and niobium plus tantalum in type 347 do not

influence the mechanical behavior of these alloys at temperatures

between 76 and 300 K.

131
The ductility of the annealed 300-series stainless steels is

generally excellent at cryogenic temperatures. Elongation and reduction of area tend to drop with decreasing temperatures, but values

generally exceed 30%.

Toughness

The austenitic stainless steels retain excellent toughness at

cryogenic temperatures. The Charpy V-notch impact toughness does

decrease with temperature, 2 but in all cases, it still far exceeds

the 20-J requirement often used in codes and standards. The Charpytoughness-vs . -temperature curves are characteristic for f.c.c. alloys

that is, ductile-to-brittle transition does not occur. The minimum

temperature for Charpy testing in the United States is 20 K. A test

facility for conducting the tests at 4 K is being used in the Soviet

TEMPERATURE. K

Fig. 1. Tensile and yield strengths of three austenitic stainless

steels—AISI types 304, 310, and 316—at temperatures

between 4 and 300 K. 2

132
Union; however, it is likely that adiabatic heating limits the usefulness of data obtained below 20 K. Low-strain-rate tests using notched

tensile specimens indicate that notch sensitivity does not develop at

A K. 2

J-integral methods have been used to measure the fracture toughness

of AISI grades 310 and 316 at temperatures to 4 K. As shown in Fig. 2,

the toughness, K^

c (J), at cryogenic temperatures exceeds the toughness

at room temperature. 1 The ratio of toughness to yield strength

[Kj

c (J)/cg

] is sufficiently high to ensure gross ductile deformation

prior to fracture.

Fatigue

Strain cycling fatigue properties at 295, 76, and 4 K were measured

for AISI grades 304L and 310 by Nachtigall8 and for grades 304L and

316 by Shepic and Schwartzberg, 7 The results for 304L, which were

essentially the same in both investigations, are shown in Fig. 3.

The fatigue resistances of 310 and 316 were superior to that of 304L,

particularly in low cycle fatigue (less than 10 4 cycles). For each

alloy, the fatigue resistance at low temperatures was superior to the

fatigue resistance at room temperature, except at the highest strain

ranges where failure occurred in less than 10 3 cycles.

The fatigue crack growth behavior of AISI grades 304, 304L, 310S,

and 316 has been determined at 295, 76, and 4 K by Tobler and Reed. 8 The data for 304L are shown in Fig. 4. Notice that the growth rates

are essentially the same at 76 and 4 K; similar behavior was observed

in each of' the other alloys. Also, notice that the growth rates at

76 and 4 K are less than those at room temperature. The best-fit lines

through the 304L data at 295 K and at 76 and 4 K form the approximate

scatter bands for the data on 304, 310S, and 316 at 295, 76, and 4 K.

Thus, it can be concluded that alloy content and temperature have minimal influence on the fatigue crack growth behavior of the 300-series

stainless steels.

Physical Properties

The physical properties of the 300-series stainless steels at

295, 77, and 4 K are summarized in Table 2. The thermal properties

have been taken from the best-fit lines used in LNG Materials and

Fluids ,

1 and the original sources are given for the elastic properties,

electrical resistivity, and magnetic permeability.

ANNEALED NITROGEN-STRENGTHENED GRADES

Nitrogen significantly increases the yield strength of austenitic

stainless steels, particularly at cryogenic temperatures. The influence

of nitrogen on the low-temperature strength and ductility of AISI 304 1 is shown in Fig. 5. Notice that the 77-K yield strength of 304N (AISI

133
Fig. 2. Fracture toughness of two austenitic stainless

steels--AISI types 310 and 316--at temperatures

between 4 and 300 K. 1

Fig. 3. Strain-cycling fatigue curves for AISI type 304L

austenitic stainless steel at 295, 76, and 4 K. b

134
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135

Ledbetter

et

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LNG

Materials

and

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lu

Efferson

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11
Fig. 4. Fatigue crack growth rates of AISI type 304L austenitic

stainless steel at 295, 76, and 4 K. 8

304 with deliberate additions of nitrogen) is approximately three times

greater than the yield strength of 304, their ultimate strengths are approximately equal, and the elongation of the 304N is superior at

temperatures below 220 K. This remarkable improvement in mechanical

properties occurs in many nitrogen-strengthened grades.

The designations and compositions of the commercially available

grades suitable for cryogenic service are summarized in Table 3. They

are divided into two basic types: the Cr-Ni-N and the Cr-Ni-Mn-N alloys.

The Cr-Ni-N alloys are 300-series stainless steels with deliberate

additions of nitrogen. In Europe, nitrogen is deliberately added to

several of the common Cr-Ni 300-series stainless steels. The equivalent AISI designations for these European grades are 304LN, 316LN,

and 347N. In United States specifications, the nitrogen range is

0.10 to 0.16% for 304N and 316N. The corresponding European

specifications permit nitrogen levels to 0.25%. The manganese in

Cr-Ni-Mn-N alloys provides austenite stability and can thus be used

to replace part of the nickel content. Manganese also increases the

solubility of nitrogen in austenite and thus permits higher nitrogen

levels; 0.4% N (max.) is typical.

136
TEMPERATURE. K

Fig. 5. Tensile properties of two austenitic stainless steels—AISI

types 304N (with 0.2% N) and 304 —at temperatures between 77

and 300 K. 12

Tensile Properties

The yield and ultimate strengths of the Cr-Ni-N 300-series stainless steels are summarized in Fig. 6. The solid lines and the broken

lines represent the data of Randak et al. 14 for 304LN and 316LN, res- pective^^ The various symbols depict the data for seven other investigations, each on one of the following alloys: 304LN, 304N, and 316LN.

As is the case for 304 and 316 (Fig. 1), alloys 304N and 304LN have

approximately 20% greater ultimate strength at cryogenic temperatures

than alloy 316LN. All three alloys have essentially the same yield

strength at and below room temperature. The nitrogen contents of the

17 heats of test materials range from 0.09 to 0.17%.

With the exception of type AISI 201, the tensile properties of

each alloy in Table 3 have been evaluated at cryogenic temperatures.

The 21Cr-6Ni-9Mn alloy has been most thoroughly evaluated. Its tensile

137
Table 3. Compositions of nitrogen-strengthened

austenitic stainless steels.

Type Nominal Composition (%)

3

Cr Ni Mn C , max N Other

AISI 201 16-18 3. 5-5.5. 5-7.

0.15 0.25, max. AISI 202 17-19 4-6 7.5-10 0. 15 0.25, max.

AISI 304N 18-20 8-10.5 2.0, max

0.08 0.10-0.16

304LN 18-20 8-20 8-12 2.0, max. 0.03 0.10-0.16

AISI 316N 16-18 10-14 2.0, max. 0.08 0.10-0.16

AISI 316LN 16-18 10-14 2.0, max

0.03 0.10-0.16

ASTM XM-10

5 19-21.5 5. 5-7.

8-10 0.08 0.15-0.40

ASTM xm-hJ 19-21.5 5. 5-7.

8-10 0.04 0.15-0.40

ASTM XM-147 17-19 5-6 14-16 0.12 0.35-0.50

ASTM XM-19^ 20.5-23.5 11.5-13.5 4-6 0.06 0.20-0.40 0.10-0.30 Cb ASTM XM-29 17-19 2.25-3.75 11.5-14.5 0.08 0.20-0.40 0.10-0.30 V

1. 5-3.0 Mo

a. For each grade: Si = 1.00, max.; P = 0.060, max.; S = 0.030, max.

b. ASTM A412

and yield strength data are summarized in Fig. 7^ 8 This alloy typically

has about 25% greater yield strength than the Cr-Ni-N alloys. Since

yield strength increases and ductility decreases with decreasing temperature, fracture resistance must be considered when selecting these

alloys for service at temperatures below 77 K.

Toughness

The nitrogen-strengthened stainless steels have excellent toughness

at room temperature. However, as shown in Fig. 8, significant toughness

losses generally occur as the temperature is reduced. 15 At 4 K, the

loss in toughness is proportional to the increase in yield strength,

as shown in Fig. 9.

Fatigue

The strain-cycling fatigue behavior of 21Cr-6Ni-9Mn has been

measured at 295, 77, and 4 K by Shepic and Schwartzberg

7 The results,

presented in Fig. 10, exhibit crossover behavior similar to that of

the 300-series stainless steels; that is, at high-cycle lives (greater

than 10 4 cycles), the fatigue strength improves with decreasing temperature, and vice versa at low-cycle lives. A comparison of the results

at 4 K (shown in Fig. 11.) indicates that the 21Cr-6Ni-9Mn alloy has

fatigue strength better than 304L but not as good as 316. 7 Thus, the

significantly higher strength of the nitrogen-strengthened grade does

not result in a comparable improvement in fatigue life.

138
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139
Fig. 8. Fracture toughness of five nitrogen-strengthened austenitic

stainless steels at temperatures between 4 and 300 K. 15

The fatigue crack growth behavior of several nitrogen-strengthened

grades has been studied at cryogenic temperatures by Read and Reed. 15

The results at 4 K depend largely on the relative austenite stability

of the alloys. For the least stable alloy evaluated, 18Cr-3Ni-13Mn,

the fatigue crack growth rates are 6 times greater at 76 K than at room temperature, and even greater at 4 K. For the most stable alloy

evaluated, 2 lCr-12Ni-5Mn , the fatigue crack growth rates fall within

the same scatter band at 295, 76, and 4 K. The fatigue behavior of

types 304N and 21Cr-6Ni-9Mn is intermediate: the growth rates are essentially the same at 295 and 76 K, but increase 2- to-3-fold at

4 K.

Physical properties

The physical properties of the 300-series nitrogen-strengthened

stainless steels are similar to those of the 300-series alloys summa- rized in Table 2). In the Cr-Ni-Mn grades, anomalies have been observed

by Ledbetter 17 in the elastic behavior and by Ledbetter and Collings 18

140
Fig. 9. Trend line for fracture toughness vs. yield strength at 4 K

for austenitic stainless steels, o — 304 (N) : A — 316 (N) — other austenitic stainless steels. 16

UJ

Fig. 10. Strain-cycling fatigue curves for 21Cr-6Ni-9Mn austenitic

stainless steel at room temperature, 76, and 4 K. 7

in some of the magnetic transitions and specific heat. In general,

the nitrogen contributes to austenite stability and thus reduces the

magnetic permeability induced in some alloys by deformation at low

temperatures

.

141
o

Fig. 11. Strain-cycling fatigue behavior of three austenitic stainless

steels—AISI types 304 and 316 and a 21Cr-6Ni-9Mn alloy—at

4 K. 7

HIGH-STRENGTH COLD-ROLLED SHEET

Most austenitic stainless steels are available in the cold-worked

condition as sheet, strip, rod, and wire. Heavier product forms, such

as plate and bar, are not usually cold-worked owing to the nonuniformity

of deformation through the thickness. Cold-rolled sheet has been extensively used at cryogenic temperatures, primarily for missile cases operating at 20 K (liquid hydrogen) or 86 K (liquid oxygen). AISI

grades 301, 302, 304L, and 310 have been evaluated for these applications . 19,20

Cold working increases the strength of stainless steels by strain

hardening and, in metastable alloys, by the strain-induced martensitic

transformation. Both of these effects are apparent in the tensile

properties of annealed stainless steels. Recall from Fig. 1 that the

ultimate strength is much greater than the yield strength as a result

of strain hardening. The differences between yield and ultimate

strength are greatest for alloys that transform to martensite during

the test; examples are AISI grades 304, 321, and 347. The strengthening of interest here is the increase in yield strength, which is

caused by the cold-working operations, not by the tensile test.

Yield strength increases are assumed to be similar for all the AISI

300-series grades.

Tensile Properties

Cold rolling causes remarkable increases in yield strength, some- what less of an increase in ultimate strength, and a severe reduction

of ductility. The influence of cold working on the yield strength

142
and elongation of AISI 310 at room temperature and 20 K is shown in

Fig. 12. 21 At 20 K, the yield strength increases 2.4-fold as the cold

work increases from 0 to 92%; the corresponding increase in ultimate

strength is 1.6-fold. The ductility at 20 K decreases linearly with

the percent of cold work, but remains 18% at 60% cold work, the extrafull-hard condition. In contrast, the room-temperature ductility is

much lower at all levels of cold work and is only 3% at 60% cold work.

Increased ductility at cryogenic temperatures is characteristic of

austenitic stainless steels and has been used to advantage in forming

operations, referred to as cryoforming . 22

Metastable alloys, such as AISI grades 301 and 302, exhibit even greater increases in yield strength as a function of cold work owing

Fig. 12. Tensile yield strength and elongation of AISI type 310

austenitic stainless steel at room temperature and 20 K as a function of percent cold work.^

143
to the strain-induced martensitic transformation. At 60% cold work,

the yield strengths at 20 K decrease with increased austenite stability

as follows: AISI 301—o = 1940 MPa, AISI 302—o = 1720 MPa, and

AISI 310o = 1600 MPaY

S The mechanical properti^l of AISI 301 in

the 60% col^-worked condition are summarized in Fig. 13. The drop in

elongation that occurs between 76 and 20 K is accompanied by an increase

in notch sensitivity and limits the usefulness of this alloy to temperatures 77 K and higher.

Toughness

Notched tensile tests are usually used to evaluate the toughness

of high-strength sheet alloys. Charpy impact tests and plane strain

fracture toughness tests have specimen-thickness requirements that

limit their applicability to sheet products. The ratio of tensile

strengths in notched and unnotched specimens is used as the index of

toughness. For stable alloys, such as AISI 310, the notched (K^, =

6 . 3)-to-unnotched tensile ratio for room and cryogenic temperatures

Fig. 13. Tensile properties of AISI type 30|1 austenitic stainless

steel, 60% cold rolled, at temperatures between 20 and

300 K. 19

144
COLD WORK, percent

Fig. 14. Notched-unnotched tensile ratio (K^ = 6.3) as a function of*

cold work for AISI type 301 austenitic stainless steel. 19

Solid line - longitudinal; broken line - transverse.

is equal to 1 ±0.1 for cold reductions to 92%.* Such is not the case

for metastable alloys. For AISI 301, the notched (K^ = 6.3)-tounnotched tensile ratio varies from 0.4 to 1.1, depending on the percent of cold work, test temperatu^^, and specimen orientation, as shown

in Fig. 14. Watson and Christian relate this behavior to the martensitic transformation. The maxima in the notched-to-unnotched tensile

ratios in Fig. 14 at 50 to 65% cold reduction were attributed to the

stabilizing influence (during cryogenic testing) of cold working up

to about 60%; cold reductions greater than 60% caused enough martensite to form during rolling to reduce notch strength at all test temperatures. The reduced toughness in the transverse specimens was attributed to the preferred orientation of the martensite rather than

to the effects of inclusions.

* K is the ratio of the maximum stress at the root of the notch to

tne nominal stress.

145
Fig. 15. Flexural fatigue (R = -1) strength of two austenitic stainless steels - annealed AISI 347 and 60% cold-rolled AISI 301

- at room temperature and 77 K. 23

Fatigue

The cryogenic fatigue behavior of AISI 301 sheet, cold rolled

approximately 60%, is shown in Fig. 15. 23 A comparison of these data

with similar data on annealed AISI 347, also obtained by Favor

et al., 23 indicates that AISI 301 with a yield strength of 1530 MPa

at 295 K has much better fatigue strength at room temperature than

annealed AISI 347 with a yield strength of 255 MPa at 295 K. At

cryogenic temperatures, however, the differences in fatigue strength

are marginal; the annealed AISI 347 has better fatigue strength at

fatigue lives less than 10 5 cycles and cold-rolled AISI 301 has

better fatigue strength at fatigue lives greater than 10 5 cycles. At

20 K, the fatigue strength of AISI 347 is superior to that of AISI

301 at all cyclic lives. Thus, in fatigue critical applications at

cryogenic temperatures, the increased yield strength of cold-rolled

301 stainless steels is of marginal value.

Notched fatigue (K^ = 3.1) tests were conducted on the same materials by Gideon et al. 24 Their results further confirm that the

high strength of cold-rolled AISI 301 stainless steel is not beneficial

in fatigue critical applications at cryogenic temperatures. The

notched fatigue behavior exhibited the same trends apparent in Fig. 15,

that is, improved fatigue strength (but to a lesser extent) at room temperature and crossover behavior at approximately 10 5 cycles at 77 K.

146
The applicability of the fatigue test results on AISI 301 to other

cold-rolled austenitic stainless steels has not been established.

However, the trends discussed should serve as a warning to those who

wish to use cold-rolled sheet in highly stressed, fatigue critical

applications: Improved tensile strength does not necessarily result

in improved fatigue strength, particularly at cryogenic temperatures.

WELDS

Welding

The austenitic stainless steels, with and without nitrogen strengthening, are readily weldable by all of the common welding processes,

providing the appropriate consumables and procedures are used. The

service experience with welded assemblies of the alloys has been

satisfactory at cryogenic temperatures. The American Welding Society

(AWS) designation system, which is similar to the AISI system for

wrought products, is generally used to classify stainless steel filler

metals. The AWS designations and chemical compositions of the filler

metals commonly used for cryogenic applications are summarized in

Table 4.

The strength of stainless steel welds at 4 K generally exceeds

the corresponding base metal strength, but toughness is usually significantly lower. Four phenomena that affect the strength and toughness

Table 4. Composition limits of selected filler metals for welding

austenitic stainless steels.

AWS Type Composition (%)

Cr Ni C, max Other

Q Covered Electrodes

E308 18-21 9-11 0.08

E308L 18-21 9-11 0.04

E310 25-28 20-22.5 0.20 0.75 Si, max. E16-8-2 14.5-16.5 7. 5-9.

5

0.10 0.50 Si, max.; 1-2 Mo

E316 17-20 11-14 0.08 2-2.5 Mo, max.

;

1-2 Mo

E316L 17-20 11-14 0.04

E330 14-17 33-37 0.25

Rod and Bare Electrod es

b

ER308 19.5-22 9-11 0.08

ER308L 19.5-22 9-11 0.03

ER310 25-28 20-22.5 0.08-0.15

ER310 18-20 11-14 0.08 2-3 Mo

ER316 18-20 11-14 0.03 2-3 Mo

a. For covered electrodes: Mn * 2.5, max. ; Si •= 0.9, max. ; P » 0.04, max. ; S = 0.03, max.

b. For rod and bare electrodes: Mn = 1.0-2. 5; Si = 0.25-60; P - 0.03 , max. ; S “ 0.03, max.

147
of the as-deposited weld metal take on added significance at cryogenic

temperatures: sensitization, ferrite content, nitrogen pickup, and

oxide inclusions. In addition, heat treatment generally has a stronger influence on the weld metal than on the base metal.

Sensitization, the grain boundary precipitation of chromium carbides,

reduces weld-metal toughness at low temperatures. 25 Consequently,

weld metals with very low carbon, such as 308L (0.04% C, max.) and

316L (0.03% C, max.) are commonly used for cryogenic applications.

Sensitization is most extreme if the weldment is given prolonged exposure

in the temperature range 500 to 800°C. In multipass weldments, the

underlying weld beads may receive sufficient exposure to these temperatures to cause significant toughness reductions in stainless steels

with greater than 0.03% C. Subsequent annealing of the welds at temperatures greater than 950°C dissolves the carbides and improves the

toughness

The ferrite content of as-deposited weld metals can be estimated

by using the Schaeffler constitution diagram. To avoid microf issuring,

the weld-metal chemistry is usually balanced to provide 4 to 10% ferrite.

However, for cryogenic service, there is concern that the ferrite,

which is a b.c.c. phase and therefore brittle at low temperatures,

may embrittle the weld metal. Szumachowski and Reid 26 have shown that

ferrite reduces toughness at 77 K for a wide range of stainless steel

weld metals. Read et al. 25 have evaluated the effect of ferrite content

on the fracture toughness of shielded metal-arc welds of AWS 316 and

316L; the results at 77 K and 4 K are shown in Fig. 16. Elmer et al. 27

have evaluated the strength and toughness of several fully austenitic

weld metals at 4 K. Their results for 20Cr-16Ni-7Mn-2.8Mo-2. IN,

17Cr-16Ni-4Mn-2.2Mo, 18Cr-20Ni-6Mn-0. 27Nb, and 16Cr-16Ni-4Mn-l . 5Mo

indicate that strength and toughness combinations comparable to the

base metal properties are possible.

Nitrogen increases the yield strength and decreases the toughness

of stainless steel weld metals. When conventional (low-nitrogen) filler

metals are used, nitrogen variations of 0.04 to 0.13% may occur, depending on the amount of air that enters the arc-shielding medium. The

greatest variations in nitrogen pickup are associated with gas metalarc (GMA) welding, where the nitrogen content ranges from 0.04 to 0.13%,

depending on the type of gas coverage, the electrode extension, and

the transfer mechanism (short-circuit is better than spray transfer).

In shielded metal-arc (SMA) welding, smaller variations are attributed

to electrode coatings; lime coatings generally give better coverage

and less nitrogen pickup than titania coatings.

The influence of oxygen content on the toughness of weld metals

at cryogenic temperatures has not yet been systematically studied.

However, it is well known that oxide inclusions form sites for the

initiation of microvoids, and that the ductile fracture resistance

148
Fig. 16. Fracture toughness of AWS 316 and 316L weld metals at 76 and

4 K as a function of ferrite content. 25

decreases as the number of microvoids increases. Some evidence on

the influence of oxide inclusions comes from a study by Whipple and

Kotecki, 28 in which a series of AWS 316L welds were produced using

three different welding processes: GMA, gas tungsten arc (GTA) , and

submerged arc (SA). The toughness at £ K increased with decreasing

inclusion content: ^SA weld—103 MPa»m^, GMA weld—161 MPa*m^, and

GTA weld—181 MPa*m 2

. Other observations in the literature r^eal the

importance of inclusions. For example, Szumachowski and Reid found

SMA deposits of 316L with a basic coating (lime coating, -15 electrode)

generally had superior toughness to those of 3161-16 deposits with a rutile coating; basic coatings are known to have lower oxygen content

and oxide inclusions than rutile coatings. 29

In summary, weld metal toughness can be improved by using lowcarbon filler metals to avoid sensitization, by balancing weld-metal

chemistry to avoid ferrite, by using low-nitrogen consumables and welding practices, and by using inert-gas-shielded welding processes or basic fluxes to reduce oxide inclusions. However, significant variations in the toughness still arise in ferrite-free welds containing

149
moderate nitrogen contents (0.004 to 0.007%), that are welded by the same process. As shown in Fig. 17, toughness varies as a function of

grain width, a feature of the weld deposit controlled by the heat input

and cooling rates. j0 To minimize these toughness variations, it is

imperative to qualify the welding materials and procedures used for

construction and to ensure that the qualified procedures are used during fabrication.

CASTINGS

Austenitic stainless steel castings are used for cryogenic equipment having heavy sections or complex geometries that make fabrication

from wrought products impractical. Typical examples are valve bodies

and piston castings for liquefaction equipment. Several large castings

have been used at temperatures down to 2 K for helium bubble chamber

Fig. 17 Fracture toughness of fully austenitic weld metals at 4 K as

a function of grain width. • —Read et al.;^ A LCP;

o — MFTF. LCP and MFTF refer to weld procedure qualification welds prepared for the Large Coil Project, Oak Ridge

National Laboratory and the Magnetic Fusion Test Facility,

Lawrence Livermore Laboratory. 30

150
bodies. The Alloy Castings Institute (ACI) designation system is generally used for stainless steel castings. The ACI types, the equivalent

AISI types for wrought products, and the chemical compositions of the

casting alloys commonly used for cryogenic applications are summarized

in Table 5.

A duplex microstructure of austenite and ferrite is typical of

most of the stainless steel casting alloys. The ferrite content of

the first six alloys listed in Table 5 may range from 0 to 40%; this

is due to chemistry variations within the allowable ranges. The other

three alloys (25Cr-20Ni, 20Cr-29Ni, and 16Cr-20Ni-10Mn) are fully

austenitic. Ferrite content influences the strength and toughness

and thus contributes to the variability observed in the mechanical

properties of the castings.

The influence of ferrite content, nitrogen, and sensitization on the toughness of stainless steel castings is similar to their influence

on weldments. Whipple and McHenry 31 evaluated the influences of 6-

ferrite and interstitial nitrogen in the strength and toughness of

stainless steel castings at 4 K using nine CF8M (nominally a 19Cr-9Ni-2Mo alloy) castings with controlled chemistries. The chromium

and nickel contents were varied to obtain two series of alloys: one series of five alloys had 0.05% N and 6-ferrite contents ranging from

0 to 28.5%, the other series of five alloys had 9 ±1% 6-ferrite and

0.02 to 0.20% N. As shown in Fig. 18, an increase in either 6-ferrite

or nitrogen content increases the yield strength and decreases ductility

at 4 K. Fracture toughness decreases with increasing 6-ferrite up to

Table 5. Compositions of austenitic stainless steel

castings.

Q Composition (%)

ACI Type Nearest

AISI Type Cr Ni C, max. Other

CF-3 304L 17-21 8-21 0.03

CF-8 304 18-21 8-11 0.08

CF-3M 316L 17-21 9-13 0.03 2-3 Mo, 1 .5 Si max. CF-8M 316 18-21 9-12 0.08 2-3 Mo, 1.5 Si max.

CF-8C 347 18-21 9-12 0.08 (8 x %C) Nb, min. or (10%C) Nb+Ta, min.

CF-20 302 18-21 8-11 0.20

CK-20 310 23-27 19-22 0.20

CN-7M — 18-22 27-31 0.07

2-

3 Mo, 1.5 Si max.

3-

4 Cu

a. For each grade: Mn = 1.5, max. ; P = 0.04, max. ; Si » 2.0, max , ; S = 0.04,

max.

151
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152
a 6-ferrite content of 17% and then remains constant (Fig. 19). The

influence of nitrogen on fracture toughness was not clearly established

because the toughness at the 9% ferrite level was already a low value,

126 to 148 MPa*m 2 at nitrogen contents of 0.02 to 0.1%. The foughness

was substantially lo^er at higher nitrogen contents: 74 MPa*m 2 at

0.

14. N and 94 MPa'm'1 at 0.20% N.

SUMMARY

The austenitic Cr-Ni stainless steels have a wide range of

mechanical properties at cryogenic temperatures. Depending upon

their chemical composition and heat treatment, the yield strengths at

4 K range from 200 to 1300 MPa. The annealed AISI 300-series grades

have moderate strength and excellent toughness at low temperatures.

Large increases in yield strength can be achieved by alloying with

nitrogen or by cold working of selected products, such as sheet.

Strength increases are accompanied by decreases in toughness and

small changes in fatigue strength. Thus, strength, toughness, and

fatigue properties should be balanced, depending on the specific

design requirements. Welds and castings can have satisfactory

properties at cryogenic temperatures, but the properties vary

considerably, depending on metallurgical variables such as ferrite

and nitrogen contents.

REFERENCES

1. LNG Materials and Fluids , D. B. Mann, ed.. National Bureau of

Standards, Boulder, Colorado (1978).

2. Handbook on Materials for Superconducting Machinery , MCIC-HB-04,

Battelle, Columbus, Ohio (1977).

3. Brickner, K. G. and Defilippi, J. D., Mechanical properties of

stainless steels at cryogenic temperatures and at room

temperature, in Handbook of Stainless Steels , McGraw-Hill,

New York (1977), pp. 20-1-20-39.

4. Materials for Cryogenic Service , Engineering properties of

austenitic stainless. International Nickel Limited, London

(1974)

.

5. Low Temperature and Cryogenic Steels , Materials Manual, U.S.

Steel Corp., Pittsburgh (1966).

6. Nachtigall, A. J., Strain cycling fatigue behavior of ten

structural metals tested in liquid helium, liquid nitrogen,

and ambient air, in: Properties of Materials for Liquefied

Natural Gas Tankage , ASTM STP 579, American Society for

Testing and Materials, Philadelphia (1975), pp. 378-396.

7. Shepic, J. A. and Schwartzberg , F. R., Fatigue testing of stainless steels, in: Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures—

I

, NBSIR 78-884,

National Bureau of Standards, Boulder, Colorado (1978),

pp. 15-68.

153
8 . Tobler, R. L. and Reed, R. P., Fatigue crack growth resistance

of structural alloys at cryogenic temperatures, in:

Advances in Cryogenic Engineering , Vol 24, Plenum Press,

New York (1978), pp. 82-90.

9. Ledbetter, H. M. , Weston, W. F. , and Naiman, E. R., Low

temperature elastic properties of four austenitic stainless

steels, J. Appl . Phys, 46:3855-3860 (1975).

10. Clark, A. F., Childs, G. E., and Wallace, G. H., Electrical

resistivity of some engineering alloys at low temperatures,

Cryogenics 10:295-305 (1970).

11. Efferson, K. R. and Leonard, W. J., Magnetic Properties of Some

Structural Materials Used in Cryogenic Applications , 0RNL-4150, Oak Ridge National Laboratory, Oak Ridge,

Tennessee (1976), p. 126.

12. Sanderson, G. P. and Llewellyn, D. T. , Mechanical properties of

standard austenitic stainless steels in the temperature

range -196 to +800C, J. Iron Steel Inst. London

207:1129-1146 (1969).

13. McHenry, H. I. and Reed, R. P., Structural alloys for

superconducting magnets in fusion energy systems, Nucl . Eng. Des. 58:219-236 (1980).

14. Randak, V. A., Wessling, W. , Bock, H. E., Steimaurer, H. and

Faust, L. , Stahl Eisen 91:1255 (1971).

15. Read, D. T. and Reed, R. P. , Toughness, fatigue crack growth and

tensile properties of three nitrogen-strengthened stainless

steels at cryogenic temperatures, in: The Metal Science of

Stainless Steel , American Society of Metals, Metals Park,

Ohio (1979), pp. 92-121.

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

selected austenitic stainless steels at cryogenic

temperatures, Cryogenics 21:415-417 (1981).

17. Ledbetter, H. M. , Anomalous low-temperature elastic behavior of

a nitrogen-strengthened chromium-manganese stainless steel.

Mater. Sci. Eng. 29:255-260 (1977).

18. Ledbetter, H. M. and Collings, E. W. , Low-temperature magnetic

and elastic-constant anomalies in three manganese stainless

steels, in: Metal Science of Stainless Steels , American

Institute of Mining, Metallurgical, and Petroleum Engineers,

New York ( 1979)

19. Watson, J. F. and Christian, J. L., Low temperature properties

of cold-rolled AISI types 301, 302, 304 ELC, and 310 stainless steel sheet, in: Low-Temperature Properties of HighStrength Aircraft and Missile Materials , ASTM STP 287,

American Society for Testing and Materials, Philadelphia

(1961), pp. 170-193.

154
20 . Watson, J. F. and Christian, J. L., Mechanical properties of

high-strength 301 stainless steel sheet at 70, -320, and

-423°F in the base metal and welded joint configuration,

in: Low-Temperature Properties of High Strength Aircraft

and Missile Materials , ASTM STP 287, American Society for

Testing and Materials, Philadelphia, (1961), pp. 136-149.

21. Christian, J. L., Gruner, J. D., and Girton, L. D., The effects

of cold rolling on the mechanical properties of type 310

stainless steel at room and cryogenic temperatures, Trans.

Am. Soc. Met. 57:199-207 (1964).

22. Schwartzberg, F. R. and Kiefer, T. F. , Properties of

Cryogenically Worked Metals , NASA contractor report No.

NASA-CR-134757 , National Aeronautics and Space Administration, Washington, D.C. (1975).

23. Favor, R. J., Gideon, D. N., Grover, H. J., Hayes, J. E., and

McClure, G. M. , Investigation of Fatigue Behavior of

Certain Alloys in the Temperature Range of Room Temperature

to -423°F , Report WADD TR 61-132, Battelle, Columbus, Ohio

(1961).

24. Gideon, D. N. , Favor, R. J., Koppenhafer, A., Grover, H. J., and

McClure, M. , Investigation of Notched Fatigue Behavior of

Certain Alloys in the Temperature Range of Room Temperature

to -423°F , ASD-TRD-62-35 1 , Battelle, Columbus, Ohio (1962).

25. Read, D. T., McHenry, H. I., Steinmeyer, P. A., and

Thomas, R. D., Jr., Metallurgical factors affecting the

toughness of 316L SMA weldments at cryogenic temperatures.

Weld. J. 59 : 104s-l 13s (1980).

26. Szumachowski, E. R. and Reid, H. F. , Cryogenic toughness of SMA

austenitic stainless steel weld metals: Part 1-Role of

ferrite, Weld. J . 57:325s-333s (1978).

27. Elmer, J. W. , McHenry, H. I., and Whipple, T. A., Strength and

toughness of fully austenitic stainless steel filler metals

at 4K, in: Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures—IV , NBSIR 81-1645,

National Bureau of Standards, Boulder, Colorado (1981),

pp. 289-302.

28. Whipple, T. A. and Kotecki, D. J., Weld process study for 316L

stainless steel weld metal for liquid helium service, in:

Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures—IV ., NBSIR 81-1645, National Bureau of

Standards, Boulder, Colorado (1981), pp. 303-321.

29. Lancaster, J. F. , The Metallurgy of Welding, Brazing and

Soldering , American Elsevier, New York (1965).

30. Whipple, T. A., McHenry, H. I., and Read, D. T., Fracture

behavior of ferrite-free stainless steel welds in liquid

helium, Weld. J. 60:72s-78s (1981).

31. Whipple, T. A. and McHenry, H. I., The mechanical properties of

stainless steel castings at 4 K, in: Austenitic Stainless

Steels at Low Temperatures , Plenum Press, New York (1982).

155
?§.;vV
AUSTENITIC-STEEL ELASTIC CONSTANTS*

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

We review recent NBS studies on austenitic-stainless-steel elastic

constants at low temperatures. By measuring velocities of longitudinal and

shear waves, we determined accurately the usual engineering elastic constants:

Young's modulus, shear modulus, bulk modulus (reciprocal compressibil ity) , and

Poisson's ratio.

Due to magnetic transitions, most of these alloys exhibit low- temperature

elastic-constant anomalies. The transitions occur over a temperature range,

but show reversible behavior and are probably second-order. A magneticsusceptibility cusp occurs near the shear-modulus transition temperature. Due

perhaps to local moments above Tp, the bulk-modulus maximum occurs at a higher

temperature. Increased nickel content lowers the transition temperature while

increased manganese raises it.

At room temperature, twenty different heats of stainless-steel 304 show

an elastic-constant variability of about 1.5 percent. For all practical

purposes these alloys exhibit elastic isotropy; texture effects, if present,

can not exceed 1.0 percent. Surprisingly, at low temperatures, elastic

constants vary several percent; this probably reflects steels in different

magnetic states.

Surprisingly, interstitial carbon and nitrogen (up to 1.3 atomic percent)

do not affect the elastic constants within 0.1 percent. We ascribe this to

the interstitial's symmetry; occupying an octahedral site, it produces an isotropic lattice distortion. With increasing manganese content, all the

elastic stiffnesses decrease linearly: the bulk modulus most and Poisson's

ratio least.

From measured elastic constants of eleven austenitic stainless steels, we deduced the effects of chromium and nickel on the elastic constants. Chromium

raises all the constants; nickel lowers them all except Poisson's ratio.

Using Kroner's averaging method, we related polycrystal 1 ine elastic

constants to monocrystall ine values. For certain cases, we describe how the

important monocrystalline elastic constants -- C,^, C^, £44 -- can be estimated from polycrystalline values.

To be published in Austenitic Steels at Low Temperatures, Eds. R. P. Reed

and T. Horiuchi, Plenum Press, New York (1983).

157
.
INTRODUCTION

We now witness many vigorous attempts to improve the properties

of austenitic stainless steels. Mechanical-property and physicalproperty improvements require two essential ingredients: an understanding based on theory and accurate measurements of important properties .

Recently, 1 the author summarized briefly studies at NBS on the

elastic constants of austenitic steels. The present paper extends

that review.

Since 1975, the elastic-properties group at NBS produced over thirty manuscripts that relate to stainless-steel physicalmechanical properties, especially the elastic constants. Experimentally, we studied the effects of several variables:

• temperature

• substitutional alloying (Cr, Ni, Mn)

• interstitial alloying (C, N)

• magnetic field

• texture (castings, welds)

• variability, sample-to-sample

Theoretically, several properties interest us:

• monocrystal/polycrystal elastic constants

• monocrystal elastic-constant prediction

• texture effects (Voigt, Reuss, Krbner)

159
• instability/elastic-constant interrelationships

• fee Fe elastic constants, monocrystal and polycrystal

• Debye temperature

• second-order phase-transition thermodynamics

• magnetostriction

• theory of 3d metals

VARIABILITY

The variability of elastic constants in 304 stainless steel was determined by measuring longitudinal and transverse ultrasonic velocities in twenty samples acquired randomly. 2 Three kinds of variations — sample-to-sample , directional within a sample, and repeated

measurements on a single sample — were reported for four elastic

constants: the bulk modulus. Young’s modulus, shear modulus, and

Poisson's ratio, denoted B, E, G, and v. Because of surprisingly

small variations, one percent or less, the principal problem became

measurement sensitivity and reproducibility. To overcome this

problem, a high-resolution measurement system was devised using

general-purpose equipment augmented with a very simple impedancetransforming amplifier and an FET transmission gate.

Effects due to frequency and directionality were negligible.

Velocity variations were quite small, 0.3 percent for v^ and 0.6

percent for v

fc , the longitudinal and transverse sound velocities.

Coupled with a 0.6 percent density variation, this leads to elasticconstant variations of 0.9 percent for E, 1.1 percent for G, and 1.4

percent for both B and v. All these are much smaller than variations implied by elastic constants reported in the literature for

these alloys. Thus, reported variabilities in these elastic con- stants must arise from experimental measurement errors, not from the

material itself. Texture effects, if present, can not exceed one percent. A similar study involving smaller numbers of specimens

gave similar results for 310 and 316 stainless steels. 3 We ascribe

the small variability to the thermal-mechanical treatment: hightemperature mechanical deformation and recrystallization.

We attribute the larger variations in B and v to small differences in the magnetic state, even above the Neel temperature. The

magnetic state affects the volume, which in turn affects the elastic

constants, especially B. More than any other elastic constant, the

Poisson ratio, v, depends on the nature of the interatomic bonding.

Several different magnetic states may occur in these alloys; those

suggested include paramagnetic, f eromagnetic , antiferromagnetic,

mictomagnetic , spin-glass, and even f errimagnetic . Thus, details of

chemical composition and thermal-mechanical treatment may affect the

magnetic state, the interatomic bonding, and Poisson’s ratio.

160
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ffi

£

(/) Z

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a

WEIGHT % Mn

Fig. 1. Compositional variation of elastic constants of Fe-Cr-Ni

alloys containing Mn. Units of E, B, and G are 10 11 N/m2

. v is dimensionless.

ALLOYING

(a) Manganese

Alloying Mn into a 304-type stainless steel decreases all the

elastic-stiffness constants and also decreases Poisson's ratio, as shown in Fig. 1. The bulk modulus decreases most, reflecting the

high compressibility of Mn atoms. Poisson's ratio and the B/G ratio

both vary slowly with Mn concentration. This indicates that Mn does

not change the nature of the chemical bonding. Manganese's unusual

161
bonding characteristics lead one to expect possible changes. By

factoring into force constants, the relationship of many-body forces

to elastic constants and bonding becomes more clear. Observation

shows that B/G = 2.005 ± 0.014 for all ten alloys. Thus, the marked

departure from a two-body-force value of 1.76 shows that many-body

forces occur in these alloys. (That many-body forces exist in

transition-metal alloys is already well known.) The surprising

result is that the ratio is nearly constant for all alloys, indicating no significant change in the character of the chemical

bonding. Poisson's ratio is nearly invariant with Mn content: v = 0.286 ± 0.001, but also different from the v = 0.261 predicted for

longitudinal-f orce-constant-only bonding.

(b) Chromium and nickel

By accurately measuring the elastic constants of eleven austenitic stainless steels of known chemical composition, we could, by

linear-least-squares analysis, determine the separate effects of Cr

and Ni on the elastic constants. 5 We found that Cr increases all

the elastic constants, including the Poisson ratio and the B/G

ratio. On the other hand, Ni decreases all the elastic stiffnesses,

but increases both v and B/G. We now try to understand this behavior from the viewpoint of 3d-metal alloy theory.

(c) Carbon and nitrogen

Nine stainless-steel 304-type alloys were studied at room temperature. 6 Carbon-plus-nitrogen contents of these alloys ranged

from 0.067 to 0.325 weight percent (0.3 to 1.3 atomic percent).

Five elastic constants — longitudinal modulus, Young's modulus,

shear modulus, bulk modulus, and Poisson's ratio — were determined

by the pulse-echo ultrasonic method. Within an experimental uncertainty of about 0.1 percent, all the elastic constants remain unaffected by C and N additions. Interstitial atoms, especially those

in isotropic, octahedral positions, should affect B more than G because they cannot move under a dilatational force. Yet, among the

nine alloys, B and G show the same variations from the mean. Present results contrast sharply with observations reported for ferritic Fe. Carbon added to alpha Fe reduces elastic stiffnesses such as

E and G by about five percent per atomic percent C! Indeed, C is

the most effective known alloying element in altering alpha Fe's

elastic constants. The essential distinction between the ferritic

and austenitic cases is the crystal structure. Body-centered-cubic

crystals are less accommodating to interstitials than are facecentered-cubic crystals. Carbon alloyed into bcc Fe occupies octahedral interstices that have tetragonal symmetry, which produce a strong tetragonal distortion; no similar distortion occurs in fee

Fe, where C occupies large octahedral interstices and the lattice

distortion is isotropic.

162
Fig. 2. Temperature variation of elastic constants of

Fe-2 lCr-6Ni-9Mn.

TEMPERATURE

Cooling produces dramatic, anomalous changes in the elastic

constants of austenitic steels, especially those high in Mn content

and/or low in Ni content. 7 19 Figure 2 shows these changes for a nitrogen-alloyed manganese stainless steel, Fe-21Cr-6Ni-9Mn. This

anomaly arises from a magnetic transition, reflected in the magnetic

susceptibility. The following principal results and conclusions

emerge from cooling studies on this and similar alloys:

1. During cooling from ambient to liquid-helium temperature,

these steels show anomalies in all their elastic constants. Both

the size of the anomalies and the transition temperatures increase

with higher Mn content and decrease with higher Ni and Cr contents.

High-Mn, low-Ni alloys actually show transition temperatures above

ambient

.

1

4

* 1

5

2. At the same time, the alloys undergo sharp antiferromagnetic-like transitions, giving rise to a cusp-like anomaly in the

163
TEMPERATURE (K) Fig. 3. Temperature variation of longitudinal elastic modulus and

magnetic susceptibility of Fe-2 lCr-6Ni-9Mn.

magnetic-susceptibility, y. Figure 3 shows the magneticsusceptibility/ temperature curve for Fe-2 lCr-6Ni-9Mn . In all these

alloys where the y-T curve has been determined, a cusp-like anomaly

occurs very near the center temperature of the elastic-constant

anomalies . 9 » 1 0 » 1

3. The shear-type elastic constants (G and E) and Poisson’s

ratio change sharply near the transition temperature, T . But the

dilatation-type elastic constant (B) changes less abruptly and

begins to soften at temperatures above T c , owing perhaps to local

moments above T . This emphasizes that there exists a magneticstate variable, perhaps the local magnetic moment, m, that affects

physical properties both below and above the paramagneticantiferromagnetic transition temperature.

4. The energy terms usually invoked for magnetic transitions

cannot explain the anomalous elastic-constant changes with temperature. In terms of conventional ferromagnetic or antiferromagnetic

164
ordering theory, a possible morphic effect is suggested. By conventional energy terms we mean domain, volume (magnetostriction), and

exchange

.

5. All alloys studied to date show qualitatively similar behavior. This occurs despite wide chemical-composition variations:

1-21 Ni, 5-25 Cr, and 1-26 Mn. Thus, the elastic-constant anomalies

arise as a characteristic of this class of alloys; they do not

depend on a particular element such as Mn.

6. Manganese affects the transition temperature more than Ni or Cr. Preliminary results for AT /Ax are +11, -3, and -2, respectively, for these three elements, where T^_ denotes transition temperature and x denotes weight-percent chemical composition.

7. Extrapolations to unalloyed fee Fe predict that it shows

antiferromagnetism at low temperatures. For Fe, the preliminary

transition temperature is 80 K.

8. All elastic-constant/temperature curves show complete re- versibility upon cooling and reheating. This suggests a second-order

transition, which is consistent with the Landau-Lif shitz 32 proof

that any magnetic phase transition can be second-order. Also, the

bulk modulus decreases during cooling through the transition, as required by second-order-phase-transition thermodynamics.

9. Bulk-modulus softening during cooling is also predicted by

quantum-mechanical models. Moruzzi et al. 33 showed for Fe, Cr, and Mn

that the effect of magnetic interactions is to increase the volume

and decrease the bulk modulus. Although these calculations were

done for ferromagnetic spin polarizations, similar repulsive forces

should exist in antiferromagnetics. The only requirement is the

existence of a magnetic moment over a scale sufficient to establish

a band structure (a few atoms).

MONOCRYSTAL/POLYCRYSTAL RELATIONSHIPS

Especially for cubic crystal symmetry, the problem of relating

the single-crystal elastic constants (C^, C , C ) to the quasiisotropic polycrystalline elastic constants (B, Cj has been studied

extensively, both theoretically and experimentally. The bulk modulus, B, because it is a rotational invariant of the elastic-stiffness

tensor, is always given by

B = (C u + 2C

12 )/3 (1)

even for a highly textured material. The expression for the shear

modulus, G, depends on the assumptions of the model. Considering

165
the self-consistent problem of an anisotropic single-crystal inclusion embedded in a corresponding isotropic matrix, and assuming

stress equilibrium at the interface. Kroner 34 showed that

G

3 + aG2 + 6G + y

= 0 (2)

where a = (5C

11 * 4C12 )/8 (3)

e = -C

44

(7C

11 - 4C12 )/8 (4)

and Y

= -C

44

(C

11 ' C12 )(C11 + 2C12 )/8 (5)

In the forward case, these relationships permit the calculation of B

and G from the C..’s, the isotropic from the anisotropic, two elastic constants froil three.

With one additional piece of information, from either theory or observation, some single relationship among the C..'s, the reverse

case can be computed also. The C „ ' s can be computed from B and G.

For a relationship among the C..'s of 316 stainless steel, we consider the existing C.. measurements on fee Fe-Cr-Ni alloys. At

least two approximate relationships emerge. First, the elastic

anisotropy is relatively constant, being 3.51 ± 0.19. Second, the y

= C /C^ ratio is relatively constant, being 0.642 ± 0.037.

Ledbetter showed that a slightly better y value results if values

for fee Fe and Ni are also included with these four alloys, the

value being y

= 0.635 ± 0.031.

For B and G values, we use those reported by Ledbetter:

1.575 ± 0.014 and G = 0.752 ± 0.015-10 11 N/m .

.3 B =

Thus, for 316 stainless steel, using y

= 0.635, the results are

C

1

= 2.06, C = 1.33, = 1.19*10 N/m . These results are consistent with most existing C.. measurements on Fe-Cr-Ni alloys.

We believe they provide the besl'' available estimate of the monocrystalline 316-stainless-steel elastic constants. The constants serve many purposes, including dislocation calculations and the ultrasonic

nondestructive evaluation of welds. Further details on this calculation occur elsewhere. 20 ’ 21

TEXTURE

Potentially, texture occurs in all polycrystalline aggregates.

By texture we mean a nonrandom distribution of crystallite orientations. Texture arises from many sources, including: mechanical

deformation, annealing (recrystallization), solidification, electrodeposition, vapor deposition (epitaxy), phase transformation, and

166
welding. For 316 stainless steel, we studied texture in welds and

in castings . 22 ’ 23

First, we describe some theoretical studies , 24-26 principally

the case of a 316-stainless-steel weld exhibiting (100) rod texture.

One expects that cubic-symmetry crystallites in welds occur either

with random orientations or with a single preferred axis parallel to

the principal heat-flow direction, denoted x^. In the X

j

-X

2 plane,

crystallites are expected to orient themselves randomly because no preferred heat-flow direction occurs in that plane. Thus, an effective fiber texture occurs with x^ being the fiber axis. Macroscopically, such a material exhibits transverse-isotropic symmetry,

with five independent elastic constants. This is equivalent to hexagonal symmetry. In Voigt's contracted C.. notation the symmetrical

elastic-stiffness matrix is:

cH

. =

1J

,H

'11

,H

'12

-,H '13

C

H

cH

0

1 1 13

C

3 3

0

'44

'44

,H

'66

( 6 )

, H

where C,,

press tne

1 H H

= ~ C^)* Thus, the theoretical problem is to exfive nexagonal-symmetry elastic constants, C„ , in terms

of the three cubic-symmetry elastic constants: C^, Cj. 2

* ^ 44

*

Present theory does not permit calculation of the elastic con- stants of textured polycrystalline aggregates of cubic-symmetry crystallites. This is a special and more difficult case of determining

the average elastic constants of a random polycrystalline aggregate.

For this latter problem, Landau and Lifshitz 35 remind us that "there

is... no general relationship between the moduli of elasticity of a polycrystal and those of a single crystal of the same substance."

The fiber-texture case involves a two-dimensional, rather than a three-dimensional, average.

However, elastic-constant bounds can be calculated for textured

aggregates. In this study we calculate the Voigt (constant strain)

167
and Reuss (constant stress) bounds and use Hill's 36 suggestion of an arithmetic average to estimate the elastic constants. This approach

turns out to be quite reasonable because the Voigt and Reuss bounds

do not differ dramatically (a maximum of 37 percent). And for the

important <10g>-te^ture case, the bounds coincide for three elastic

constants: Cj

3

, ^33’ an<* ^44’ Thus, we ^now these three constants

unambiguously. Kroner and wawra 37 explain the systematics of this

bounds coincidence, and they give the following expressions for the

r" ' e • VS!

’ll Cj + 30C3 ? (3C11

+ C

12 + 2C

4A )

where

,H

'12

,H

'13

,H

'33

,H

'44

,H

'66

C

2 + 6C3

C + 8BC

3

?( Cn

+ 3C

12

" 2CW

C

2 “ 4eC

3 - c

12

11

^(C

1 - C2

) - 46C

3 '44

5

(c n

- c

i 2

> “ J

(cn

* c

i2

+ 2c

aa>

(7)

5C, 3C

11 + 2C12

+ 4C

44

5C. C

11 + 4C12 + 2C

44 ( 8 )

C

11 " C12 ” 2C

44

where B * 1/20 for <100> fiber texture.

Apparently, relationships for the S 's, which give the Reuss

bounds by matrix inversion, do not occur in the literature. These

are

5

11 * §

(6S

11

+ 2S

12

+ S A4>

5

12

= 8^ 2S

11 + 6S

12

" S 4A^

168
S ( 9 ) 13 “ S

]

2

S

33

= S

11

S

44

“ S 44

S 66

= 2(S

11 " S12 ) =

f (2S H

" 2S

12 + S 44^

In^deriving these relationships, one must remember that the invers|

3 ^ ourt^“ ran^ elasticity tensor relates to the invers|

or

J its associated 6x6 matrix by a factor multiplied times C...

This factor is 4 if both a and 6 exceed 3, is 2 if either a or^B

exceeds 3, and is 1 if neither a nor B exceeds 3.

Fig. 4. Angular variation of Young's modulus for 316 stainless

steel for <001> texture. Vertical direction is x^ (hex),

about which figure has complete rotational symmetry.

Average denotes simple arithmetic average of Voigt and

Keuss bounds, which correspond to constant strain and

constant stress, respectively.

169
Fig. 5. Angular variation of torsional modulus for 316 stainless

steel for <001> texture. See figure caption 4.

Table 1 shows the predicted single-crystal elastic constants of

316 stainless steel. Other principal results of the study—elasticstiffness, elastic compliance, and sound velocity—occur in Table 1,

which shows Voigt and Reuss bounds together with their arithmetic

average for the <100> ideal fiber texture. Tjo obtain sound velocities, we used a mass density of 7.958 g/cm . From these results,

the polar plots shown in Figs. 4-6 were determined using well-known

relationships. These figures give the angular variation of Young's

modulus, E, torsional modulus, T, and sound velocities, both longitudinal, v , and transverse, v . All curves in Fig. 6 represent

arithmetic averages of Voigt and Reuss bounds.

These calculations show the following principal results:

1. Texture induces large elastic-constant changes (departures

from isotropic values).

H H

2. Shear elastic cgnstants^ and C^, change more than

longitudinal^constan^s , and , and more than the off-diagonal

constants, and

3. C

44

offers the best opportunity for detecting texture and

estimating its amount. C

44 corresponds to a shear wave propagated

along the fiber axis (x^) and vibrating in the x

i “x

2 plane, which is

elastically isotropic.

170
Table 1. Elastic stiffness, elastic compliance, and sound

velocities for 316 stainless steel with <100>

fiber texture.

Voigt Reuss Average Isotropic

cH

11

2.473 2.265 2.369 2.577

C

H

L

12

0.926 1.133 1.029 1.073

C

H

i3

1.326 1.326 1.326 1.073

C

H

33

2.072 2.072 2.072 2.577

C

H

L

44

1.174 1.174 1.174 0.752

C

H

L

66

0.774 0.566 0.670 0.752

S Hb ll

0.617 0.736 0.676 0.514

SH

b

12

-0.029 -0.148 -0.088 -0.151

S H

13

-0.376 -0.376 -0.376 -1.515

S H

33

0.964 0.964 0.964 0.514

S Hb 44

0.852 0.852 0.852 1.330

S Hb

66

1.293 1.766 1.530 1.330

\ ( V 0.557 0.534 0.546 0.569

\

(X

3

) 0.510 0.510 0.510 0.569

v

t (Xj) 0.312 0.267 0.289 0.307

V

t (X3> 0.384 0.384 0.384 0.307

aTT . Units on C . . , S . . , and

ij iJ

in ii M/ 2 v are 10 N/m , 10 ^m^/N, and cm/ys

,

respectively.

171
0.6

SOUND VELOCITY, cm/fis

Fig. 6. Angular variation of three sound velocities. See figure

caption 4.

^

4 . If high attenuation precludes accurate measurement,

provides a good lower-attenuation alternative. corresponds

to

J a longitudinal wave propagating along the fiber axis.

5. For the principal crystallographic directions, with one exception, if texture increases v^ it decreases v , and vice versa. Hardly surprising, this means simply that stiffening in a particular

direction corresponds to softening in transverse directions.

6 . ,H and

e difrepreH

Elastic constants affected least by ^exture are C^,

S^. This result^ at firgt surprising since depends^on th

ference between and arises presumably because

sents a shear wave both propagated and polarized in the x

^ -x2 plane,

an elastically isotropic plane with properties determined, for all

textures, by averaging over all directions perpendicular to the

fiber axis.

7. Finally, we remark concerning the bulk modulus of the

f iber-textured^aggregate. It equals identically the cubic-symmetry

bulk modulus, + 2C.

2 )* It is a scalar. Hydrostatic stress

would dilate uniformly tne textured aggregate. This contrasts

sharply with a microscopically hexagonal material that behaves anisotropically, altering its axial length ratio, in response to hydrostatic stress. This means that welds of cubic-symmetry alloys do

172
not develop pressure-induced internal stresses due to texture. Such

stresses would occur, of course, in lower- symmetry-alloy welds.

Another study26 extends the theory to include [110] and [111] rod

textures

.

A general conclusion of these studies is that sound velocity

provides a powerful probe for detecting, characterizing, and, at

least in simple cases, quantifying texture.

INSTABILITY

Many austenitic steels exhibit crystal-structure instability

when subjected to changes of temperature, stress, or pressure.

Above, we described the temperature-induced magnetic transitions in

these alloys. Without these magnetic transitions, which tend to

stabilize the crystal structure, many more crystal-structure instabilities might occur. Thus, magnetic and structural transitions are related; both relate to the elastic constants.

0 0.2 0.4 0.6 0.8 1.0

Fig. 7. Reduced monocrystalline elastic stiffnesses for fee

elements

.

173
Recently, Ledbetter 27 conceived a new relationship between

monocrystal elastic constants and structural instability. In 1976,

Ledbetter and Moment 28 published Fig. 7, a diagram used to describe

the unusual elastic behavior of fee Pu. This figure contains con- siderable information. For the usual three independent cubicsymmetry Voigt elastic stiffnesses—C^, C^2* ^44—it plots two

dimensionless ratios: y = C^/C. , and x = C /C . The Cauchy relationship, C^„ = C, ,, is a straight line from (0,0) to (1,1). The

point (0.5, 0.5) identifies the prediction of an fee nearestneighbor-only central-force model. Most fee elements occur in the

upper half of the diagram and mostly in the left half. The boundaries x = 0 and y

= 1 correspond to mechanical-stability conditions:

C,>0 and (C^. - C^) 5^. ^he third mechanical-stability condition,

(C n + 2C^2 ) > u, or y>-0.5, does not appear in the diagram; but this

is of no consequence because no fee elements exhibit negative

values. Straight negative-slope lines emanating from (0,1) correspond to various values of Zener's anisotropy ratio

A - 2C

44 /(C11 - C12 ) <10)

All points in the diagram's field correspond to mechanically stable

crystal structures.

That Th and Pu occur to the right of x = 0.5 provides a crucial

basis for this argument. All elements to the left of x = 0.5 are stable; they exhibit a single crystal structure between T = 0 and

the melting point. Both Th and Pu are unstable. At high temperatures, Th transforms to a body-centered-cubic crystal structure.

Plutonium is notoriously unstable. Between ambient and melting temperatures it exhibits six allotropes: a (monoclinic), 8 (monoclinic), y (orthorhombic), 6 (fee), 6' (tetragonal), e (bcc) . (One can conjecture that lower temperature may reveal further allotropes,

perhaps triclinic.) Thus, despite high values of x, representing

high elastic rigidity with respect to ( 100) [001 ]-type deformation,

both Th and, especially, Pu are unstable. It may be significant

that elements on the diagram's left tend to be dominated by s-p

electrons in their cohesive properties. In Th and Pu the d and f

electrons enter strongly into cohesion. The tendency of Th and Pu

toward chemical instability, despite their mechanical stability, en- courages one to examine other systems, especially alloys.

Figure 8 shows the variation of x with Ni content for Fe-Ni

alloys. Of course, this composition region is well known for its

instability and the incipient fcc-bcc transformation. That x de- creases before the structural transformation occurs suggests a possible magnetic transition that stabilizes the fee structure for

awhile. Magnetic-property measurements between 25 and 35 percent Ni

would be very useful in understanding this unusual pretransformation

behavior.

174
Fig. 8. Variation of x = C, , /C. , with Ni content for Fe-Ni

11 44 11

alloys

.

Figure 9 shows a y-x diagram for austenitic (fee) stainless

steels, which consist mainly of Fe, Cr, Ni, and Mn. These alloys

are well known for their instability. At low temperatures they

undergo magnetic and/or structural (to bcc) transformations. Plastic deformation of these alloys induces transformations to either

bcc or eph (close-packed-hexagonal) structures. This diagram

includes points for fee Fe, Co, and Ni. Iron’s allotropy is well

known: bcc at low and high temperatures, fee between. Cobalt, fee

at high temperatures transforms to eph at 603 K. Elastic constants

of hypothetical fee Mn and hypothetical fee Cr remain, apparently,

unestimated. (Recently, Moruzzi et al. 33 estimated the bulk modulus

of fee Mn to be 2.91 Mbar, 95 percent of that for bcc Fe. The bulk

modulus equals (C^ + 2 C )/3. Thus, unless the C of fee Mn

behaves abnormally, its elastic constants may resemble Fe's.)

A possible explanation of these instabilities lies in the force

constants, which relate more directly to forces between pairs of

atoms in solids. Details of a force-constant model occur elsewhere . 27

BINARY ALLOYS: Fe-Ni

From a theoretical viewpoint, austenitic steels contain considerable complexity. The four principal constituents—Fe, Cr, Ni,

175
Fig. 9. Reduced monocrystalline elastic stiffnesses for austenitic

steels

Mn—are all 3d transition metals and all are magnetic. Only re- cently has theory succeeded in predicting the physical properties of

these elements and their binary alloys. Thus, to better understand

the properties of Fe-Cr-Ni-Mn alloys, we should focus also on simpler binary systems; iron-nickel, for example. Ledbetter and

Reed 29 reviewed the elastic constants of these alloys. Ledbetter,

Naimon, and Weston 30 measured the low-temperature elastic constants

of Fe-30Ni, invar. For these alloys. Fig. 10 31 shows the compositional variation of the Poisson ratio. Within the fee phase field,

v shows both a maximum and a minimum. We pointed out above that v, more than most elastic constants, depends on the nature of the

interatomic forces. Figure 10 suggests that with increasing Fe,

before the fcc-bcc phase transition, the magnetic state undergoes a change, which stabilizes the fee. Instead of increasing toward the

Born mechanical-instability limit, v turns downward to a value near

0.22. It remains unclear whether the increase in \) with further Fe

alloying arises because of another magnetic-state change or from the

natural result of alloying effects.

ACKNOWLEDGMENT

This study arose mainly from research sponsored by DARPA and by

the U.S. Department of Energy, Office of Fusion Energy. M. W.

Austin of NBS provided much valuable assistance, in particular,

highly skilled experimentation.

176
Fig. 10. For Fe-Ni alloys, compositional variation of Poisson

ratio.

REFERENCES

1. H.

2. H.

3. H.

4. H.

5. H.

6. H.

7. H.

8. H.

M. Ledbetter, Stainless-steel elastic constants at low

temperatures: a review, j.n "Advances in Cryogenic Engineering," Vol. 30, Plenum, New York (forthcoming).

M. Ledbetter, N. V. Frederick, and M. W. Austin, Elasticconstant variability in stainless steel 304, J. Appl Phys.

51:305 (1980).

M. Ledbetter, Sound velocities and elastic constants of

steels 304, 310, and 316, Metal Sci. 14:595 (1980).

M. Ledbetter, Manganese contributions to the elastic con- stants of f ace-centered-cubic Fe-Cr-Ni stainless steels,

submitted for publication.

M. Ledbetter and M. W. Austin, Elastic constants of polycrystalline austenitic stainless steels, in preparation.

M. Ledbetter, Effects of carbon and nitrogen on the elastic constants of stainless-steel 304, submitted for

publication.

M. Ledbetter, W. F. Weston, and E. R. Naimon, Lowtemperature elastic properties of four austenitic stainless steels, J. Appl. Phys 46:3855 (1975).

M. Ledbetter, Anomalous low-temperature elastic behavior

of a nitrogen-strengthened chromium-manganese stainless

steel. Mater. Sci. Eng. 29:255 (1977).

177
9 .

10 .

11 .

12 .

13.

14.

15.

16.

17.

18.

19.

20 .

21 .

22 .

23.

24.

25.

26.

27 .

H. M. Ledbetter and E. W. Collings, Low-temperature magnetic

and elastic-constant anomalies in three manganese stainless steels, in "Metal Science of Stainless Steels," AIME,

New York (1980)

E. W. Collings and H. M. Ledbetter, Sound velocity anomalies

near the spin-glass transition in an austenitic stainless

steel alloy, Phys. Lett. 72A:53 (1979).

H. M. Ledbetter, Anomalous low-temperature elastic-constant

behavior in Fe-13Cr-19Mn, Metall. Trans. 11A:543 (1980).

H. M. Ledbetter, Room-temperature elastic constants and

low-temperature sound velocities for six nitrogenstrengthened stainless steels, Metall. Trans. 1 1 A : 1 06

(1980).

H. M. Ledbetter, Stainless steel elastic constants at low

temperatures, J. Appl. Phys. 52:1587 (1981).

H. M. Ledbetter, Low-temperature elastic-constant anomalies

in Fe-Cr-Ni-Mn alloys, Physica (forthcoming).

H. M. Ledbetter and M. W. Austin, Elastic constants of an Fe-5Cr-25Mn austenitic steel, 76-400 K, submitted for publication . H. M. Ledbetter and M. W. Austin, Anomalous low-temperature

elastic-constant behavior of Fe-20Cr-16Ni-6Mn, submitted

for publication.

H. M. Ledbetter and E. W. Collings, Elastic-constant anom- alies at the N£el transition in Fe-18Cr-3Ni-12Mn, submitted for publication.

H. M. Ledbetter, Low-temperature variability of stainlesssteel-304 elastic constants, submitted for publication.

H. M. Ledbetter, Low-temperature elastic constants of a man- ganese-alloyed stainless steel, in preparation.

H. M. Ledbetter, Predicted single-crystal elastic constants

of stainless-steel 316, Brit. J. Non-destr. Test. 34:286

(1982)

H. M. Ledbetter, Predicted single-crystal elastic constants

of 304 stainless-steel, submitted for publication.

H. M. Ledbetter and M. W. Austin, Anisotropic elastic properties of 316L welds, in preparation.

H. M. Ledbetter and M. W. Austin, Elastic constants and internal friction of some stainless-steel castings, in

preparation

H. M. Ledbetter, Single-crystal elastic constants in non- destructive evaluation of welds, _in "Review of Progress

in Quantitative Nondestructive Evaluation," Plenum, New

York (1982), pp. 619-624.

H. M. Ledbetter, Single-crystal elastic constants in ultrasonic study of welds, submitted for publication.

H. M. Ledbetter, Sound velocity as a texture probe: application to stainless-steel welds, submitted for publication.

H. M. Ledbetter, Elastic constants and instability in facecentered-cubic crystals, submitted for publication.

178
28 . H. M. Ledbetter and R. L. Moment, Elastic constants of facecentered-cubic plutonium, Acta Metall. 24:891 (1976).

29 H. M. Ledbetter and R. P. Reed, Elastic properties of metals

and alloys, I. iron, nickel, and iron-nickel alloys, J.

Phys. Chem. Ref. Data 2:531 (1973).

30. H. M. Ledbetter, E. R. Naimon, and W. F. Weston, Lowtemperature elastic properties of invar, in "Advances in

Cryogenic Engineering," Vol. 22, Plenum, New York (1977),

pp. 174-181.

31. H. M. Ledbetter, Single-crystal and calculated polycrystalline elastic constants of iron-nickel alloys, a review, in

preparation.

32. L. D. Landau and E. M. Lifshitz, "Statistical Physics,"

Pergamon, London (1959).

33. V. L. Moruzzi, J. F. Janak, and A. R. Williams, "Calculated

Electronic Properties of Metals," Pergamon, New York

(1978).

34. E. KrSner, Berechnung der elastischen Konstanten des

Vielkristalls aus den Konstanten des Einkristalls, Z

.

Phys . 151:504 (1958).

35. L. D. Landau and E. M. Lifshitz, "Theory of Elasticity,"

Pergamon, London (1959).

36. R. Hill, The elastic behaviour of a crystalline aggregate,

Proc. Phys. Soc . Ab5:349 (1952).

37. E. Krtiner and H. H. Wawra, Coincidence of bounds for some elastic moduli of polycrystals with fibre texture,

Philos. Mag. A 38:433 (1978).

179
S&v
Elastic constants of an Fe-5Cr-26Mn austenitic steel, 76 - 400 K

H. M. Ledbetter and M. W. Austin

Fracture and Deformation Division, National Bureau of Standards, Boulder,

Colorado 80303, USA

Abstract

By measuring longitudinal-mode and transverse-mode sound velocities at

frequencies near 10 MHz, we determined the complete engineering elastic

constants—bulk modulus, shear modulus. Young modulus, Poisson ratio—for an Fe-5Cr-26Mn austenitic steel between 76 and 400 K. Due to a magnetic transition, all elastic constants behave anomalously below about 360 K. The bulk

modulus begins to soften during cooling at some higher temperature. Except

for Poisson's ratio, below the 360-K magnetic transition, all elastic constants

resume an apparently normal temperature dependence. After increasing abruptly

at the magnetic transition, Poisson's ratio increases with decreasing temperature 1

Key Words: Austenitic steel; bulk modulus; elastic constants; iron alloys;

low temperatures; manganese alloys; mechanical properties; physical

properties; Poisson ratio; shear modulus; Young modulus.

181

Introduction

We now witness much vigorous austenitic-steel research' [1,2]. Many new,

interesting alloys transcend the limitations of the conventional 304-type

stainless steel, Fe-18Cr-8Ni.

Notably, Mn-alloyed austenitic steels provide attractive physicalmechanical properties combined with lower cost due to Ni-content reduction.

Although currently much studied, Mn-alloyed austenitic steels we-re first

considered carefully in Germany in the 1930's [3].

Alloying Mn into fee Fe must cause drastic physical-property changes.

Manganese's peculiar elemental properties—a 58-atom unit cell, a bulk modulus

0.35 that of Fe's, antiferromagnetic electronic interactions, negative lowtemperature thermal expansivity—preclude any reliable predictions of such

effects. Despite the proximity of Mn and Fe in the first long row of the

periodic table of elements, their properties differ dramatically. From

available elastic constants of fee Ni and Mn, Vegard's law predicts that Mn

lowers all the elastic stiffnesses and Poisson's ratio, v: B by 0.33 percent,

E by 0.18 percent, G by 0.18 percent, and v by 0.16 percent per percent

solute. Here, we assume that the unknown elastic constants of fee Fe do not

differ significantly from those of fee Mi; because Fe's atomic volume exceeds

Ni's by eight percent, fee Fe may be slightly softer elastically.

One may ask also whether Mn affects chemical bonding in Fe-Cr-Ni alloys.

As described above, Mn is truly a maverick metallic element. In considering

chemical bonding in both nonmetals and metals, Pauling [4] referred to "the

anomalous Mn radius" and to Mn's "striking abnormality in behavior." We know

that elastic stiffness depends relatively sensitively on atomic radius, r, -4

varying as r . Thus, Mn's larger atomic volume arising from magnetic repulsions

explains, in part, its lower elastic stiffness.

183
Previously, we reported the low-temperature elastic constants of an

experimental Soviet austenitic steel, Fe-13Cr-19Mn [5]. Owing to a magnetic

transition near 274 K, this steel showed anomalies in all elastic constants.

Below the magnetic transition, all constants resumed apparently normal temperature dependence. From other studies [6], we learned that higher Ni content

lowers the transition temperature and that higher Mn content raises it.

The present study focused on a high-Mn-con tent , low-Ni-content alloy,

Fe-5Cr-26Mn. Thus, larger anomalies were expected at a higher transition

temperature. Especially for cryogenic applications, this alloy offers many

desirable properties: stable austenite phase; good strength, ductility, and

toughness; relatively low thermal expansivity [7].

Material

Samples were obtained from a plate produced in a 45,000-kg commercial

pilot electric-furnace heat. Bottom-poured 14,000-kg ingots were reheated to

12.00°C and rolled into 4.-0-cm plates. These were solution treated at 1050°C

for 1.5 h and water quenched. Chemical composition (weight percent ladle

analysis) was: 25.6 Mn, 4.67 Cr, 0.97 Ni, 0.86 Si, 0.22 C, 0.043 Nb , 0.040 N,

0.024 P, 0.009 Al, 0.001 S.

Measurements

Sound velocities were determined by a method described in detail

previously [8], Briefly, 1.5-c.m cubes were prepared by grinding so that

opposite faces were flat and parallel within 5 ym. Quartz piezoelectric

crystals with fundamental resonances between 4 and 7 MHz were cemented with

phenyl salicylate to the specimens. 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

184
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. Figure 1 shows a typical

pulse-echo pattern. The sound velocity was computed by

v = 2 £ / 1 (1)

where l denotes specimen length, and t the round-trip transit time. On the

oscilloscope, t was the time between adjacent echoes, the first and second

echoes usually being measured, and within these the time between leading

cycles. Elastic constants were computed from the general relationship

C = pv

2 (2)

where p denotes mass density. The usual engineering elastic constants are

related to the longitudinal and transverse sound velocities, v

£ and v^_, by

2 longitudinal modulus = C

£ = pv

£ (3)

2 shear modulus = G = pv^ (4)

bulk modulus = B = - (4/3)G (5)

Young's modulus = E = 3GB/ (C - G) (6)

Poisson's ratio = v = (E/2G) - 1

= (1/2) (C^ - 2G)/(C

£ - G) (7)

For low-temperature measurements, we used a cryostat similar to one

described previously [9] and a cyanoacrylate specimen-transducer adhesive.

Above-ambient measurements were made in a stirred, heated mineral-oil bath.

Results

Table 1 gives for selected temperatures both sound velocities and elastic

constants. Figures 2 and 3 show the temperature variations. Maximum uncertainties are believed to be 1.0 percent for G and E and 1.5 percent for B and v.

185
The results show the following principal features: (1) near 335 K, a

transition occurs that affects all the elastic constants; (2) upon cooling

through the transition, the bulk, shear, and Young’s moduli decrease and

Poisson’s ratio increases; (3) except for Poisson’s ratio, below the transition temperature all elastic constants resume apparently normal temperature

variations; (4) at low temperatures, Poisson's ratio, v, shows a nearly

linear, but anomalously negative dv/dT; (5) above the transition temperature,

all elastic constants show usual temperature variations, except for the bulk

modulus, B, which shows a nearly zero dB/dT, which is usually negative, except

as zero temperature is approached.

Discussion

We believe the transition near 335 K is magnetic, the high-temperature

phase being paramagnetic. The magnetic state of order in the low-temperature

phase remains uncertain. But magnetic-susceptibility measurements on three

austenitic alloys—Fe-22Cr-l 3Ni-5Mn , Fe-2 lCr-6Ni-9Mn , Fe-18Cr-3Ni-13Mn

revealed a sharp "antiferromagnetic-like" cusp-like anomaly [6] due perhaps to

a spin-glass condensation [10,11]. The elastic-constant-versus-temperature

curves for those three alloys resemble strongly the present curves for Fe-5Cr-26Mn.

The magnetic transition is probably second-order. Landau and Lifshitz

showed that any magnetic transition can be second-order [12]. The bulk-modulus

decreases on cooling, as required by second-order-phase-transformation thermodynamics [ 13]

Moruzzi et al. [14] showed for Fe, Cr, and Mn that the effect of magnetic

interactions is to increase t'he volume and decrease the bulk modulus. These

effects can be quite large. Manganese shows the largest effect. Calculations

using a sophisticated density-functional theory, but a nonmagnetic model,

112 112

predict B = 2.91*10 N/m . Observation shows [15] that B = 0.597*10 N/m ,

186
only forty-three percent of that predicted! These authors considered only

ferromagnetic spin-polarization, but they "speculate that similar repulsive

forces exist also in antiferromagnets . All that is required is the existence

of a magnetic moment over a distance scale sufficient to establish a band

structure (a few atoms)... it does not matter if... the moments are disordered

or antif erromagnetically ordered instead of being in perfect ferromagnetic

alignment ... the volume expansion is independent of the spin direction" [16].

Thus, regardless of the low-temperature magnetic-order state, one expects

during cooling a volume increase and a bulk-modulus decrease.

Wohlfarth [17] discussed in detail the influence of ferromagnetism on the

bulk, shear, and Young’s moduli of alloys. The change in elastic constant, C,

at the magnetic transition has four parts:

AC = AC, + AC +. AC + AC (8)

A 0) y m

where AC, arises from linear magnetostriction, AC from volume magnetostriction,

A U) AC^ from shape effects, and AC^ from magnetization:

2 2

1 3 F M

ACm

= C(M) - C (0) = r2 3r / cn (9)

where M denotes magnetization, F the magnetic Helmholtz free energy, and w the

volume strain.

However, another important influence on C exists: the dependence of the

-4

nonmagnetic bulk modulus on volume. Gilman [18] showed that B varies as r ,

where r denotes atomic radius. Thus, quite apart from B varying with magnetostriction or with second strain derivatives of the magnetic free energy, B

depends also on volume. The volume expansion on cooling is, as described

above, a magnetic effect related closely to the usual spontaneous volume

magnetostriction

.

] 87
It is useful to compare the elastic constants of Fe-5Cr-26Mn with those

of the more-familiar 304-stainless-steel alloy [19]. Table 2 shows the

comparison as ratios at two temperatures, 0 and 400 K, in the magnetic and

nonmagnetic states, respectively. These ratios reflect a general relative

elastic softness of Fe-5Cr-26Mn. The nonmagnetic-state lowering of Poisson's

ratio is especially interesting. Foster and Franz [15] pointed out that

"Poisson's ratio depends to a much greater extent on the conditions of bonding

than do the other elastic coefficients." Thus, these two alloys may differ in

their paramagnetic-state bonding.

Acknowledgment

This study arose from a program sponsored by the U.S. Department of

Energy, Office of Fusion Energy.

References

1. E. W. Collings and H. W. King (eds.). The Metal Science of Stainless

Steels (AIME, New York, 1978).

2. T. Horiuchi and R. P. Reed (eds.). Austenitic Steels at Low Temperatures

(Plenum, New York, forthcoming).

3. G. Masing, Lehrbuch der Allgemeinen Metallkunde (Springer-Verlag, Berlin,

1950), pp. 578-580.

4. L. Pauling, The Nature of the Chemical Bond (Cornell U.P., Ithaca, 1960),

pp. 254, 419.

5. H. M. Ledbetter, Anomalous low-temperature elastic-constant behavior in

Fe-13Cr-19Mn, Metall. Trans. 1 1 A , 543-544 (1980).

6. H. M. Ledbetter and E. W. Collings, Low-temperature magnetic and elasticconstant anomalies in three manganese stainless steels, in The Metal

Science of Stainless Steels (AIME, New York, 1978), pp. 22-40.

7. H. Yoshimura, N. Yamada, H. Honma, and T. Ito, Microstructures, lowtemperature toughnesses, and thermal-expansion coefficients of highmanganese chromium austenitic steels, Trans. Iron Steel Inst. Japan 16 , 98-105 (1976).

188
8. H. M. Ledbetter, M. F. Frederick, and M. W. Austin, Elastic-constant

variability in stainless-steel 304, J. Appl. Phys. 51 , 305-309 (1980).

9. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Elastic properties of

two titanium alloys at low temperatures. Cryogenics 246-249 (1974).

10. E. W. Collings and H. M. Ledbetter, Sound-velocity anomalies near the

spin-glass transition in an austenitic stainless-steel alloy, Phys. Lett.

72A, 53-56 (1979).

11. E. W. Collings and S. C. Hart, Low temperature magnetic susceptibility

and magnetization studies of some commercial austenitic stainless steels.

Cryogenics 19_, 521-530 (1979).

12. L. D. Landau and E. M. Lifshitz, Statistical Physics (Pergamon, London,

1959), p. 445.

13. Ref. 12, p. 438.

14. V. L. Moruzzi, J. F. Janak, and J. R. Williams, Calculated Electronic

Properties of Metals (Pergamon, New York, 1978), p. 6.

15. W. Roster and H. Franz, Poisson's ratio for metals and alloys, Metall.

Rev. 6 , 1-55 (1961).

16. Ref. 14, p. 164-166.

17. E. P. Wohlfarth, The influence of ferromagnetism on the bulk, shear, and

Young's moduli of alloys, J. Phys. F: Metal Phys. 6^, L59-L62 (1976).

18. J. J. Gilman, Micromechanics of Flow in Solids (McGraw-Hill, New York,

1969), pp. 29-41.

19. H. M. Ledbetter, Stainless steel elastic constants at low temperatures,

J. Appl. Phys. .52 , 1587-1589 (1981).

List of Figures

Fig. 1. Oscilloscope display of pulse-echo pattern for Fe-5Cr-26Mn. Transit

time was measured between crests of leading cycles of consecutive

echoes

.

Fig. 2. Temperature dependence of the longitudinal and transverse sound-wave

velocities of Fe-5Cr-26Mn.

Fig. 3. Temperature dependence of four elastic constants of Fe-5Cr-26Mn:

E = Young's modulus, G = shear modulus, B = bulk modulus, v = Poisson's

ratio.

189
T

K

395

380

360

340

320

300

280

260

240

220

200

180

160

140

120

100

80

76

Table 1. Dynamic elastic properties of Fe-5Cr-26Mn steel

at selected temperatures

E G

1

H N/m2 B V

aw/ yb 1U IN/ III

0.5480 0.3136 1.923 0.765 1.316 0.256

0.5486 0.3144 1.932 0.769 1.316 0.255

0.5492 0.3151 1.939 0.773 1.317 0.255

0.5474 0.3138 1.923 0.766 1.311 0.255

0.5389 0.3034 1.817 0.716 1.305 0.268

0.5401 0.3037 1.821 0.718 1.313 0.269

0.5410 0.3040 1.826 0.719 1.318 0.269

0.5422 0.3045 1.832 0.721 1.326 0.270

0.5435 0.3048 1.837 0.723 1.334 0.270

0.5447 0.3054 1.844 0.726 1.341 0.271

0.5458 0.3058 1.851 0.728 1.348 0.271

0.5471 0.3064 1.858 0.731 1.355 0.271

0.5483 0.3070 1.865 0.733 1.361 0.272

0.5492 0.3077 1.873 0.737 1.365 0.271

0.5504 0.3082 1.880 0.739 1.372 0.272

0.5513 0.3086 1.886 0.741 1.377 0.272

0.5524 0.3089 1.890 0.743 1.384 0.272

0.5528 0.3071 1.892 0.743 1.387 0.273

190
Table 2.

Ratios of Fe-5Cr-26Mn elastic constants to those

of 304-type stainless steel.

T(K) B E G V

0 0.88 0.90 0.90 0.98

400 0.84 0.94 0.94 0.87

191
Fig. 1. Oscilloscope display of pulse-echo pattern for Fe-5Cr-26Mn. Transit

time was measured between crests of leading cycles of consecutive

echoes

192
S7//LU0

‘

TEMPERATURE, K

Fig. 2. Temperature dependence of the longitudinal and transverse soundwave velocities of Fe-5Cr-26Mn.

193

t

,

cm//j.s
SHEAR

MODULUS,

YOUNG'S

MODULUS,

10

11

N/m

2

10

11

N/m

2

TEMPERATURE, K

Fig. 3. Temperature dependence of four elastic constants of Fe-5Cr-26Mn

E = Young's modulus, G = shear modulus, B = bulk modulus, v

= Poisson's ratio.

194

POISSON'S

RATIO

BULK

MODULUS,

10

11

N/m

2
WELDMENTS AHV CASTINGS

195

WELDMENTS AND CASTINGS PROGRAM

LEADER: H. I. McHenry, NBS

STAFF: Y.-W. Cheng, H. M. Ledbetter, NBS

OBJECTIVES:

(1) To investigate the metallurgical factors that affect the mechanical

properties of stainless-steel weldments and castings at cryogenic

temperatures.

(2) To contribute to the development of improved filler matals for

welding stainless steels for liquid-helium service.

(3) To evaluate the mechanical properties of weldments and castings at

4 K. (4) To develop methods of detecting defects and evaluating their

significance in stainless-steel weldments.

RESEARCH PAST YEAR (1982)

(1) The five CF8M stainless-steel castings with varying nitrogen

contents and constant delta-ferrite content were tested at 4 K. (2) The metal lographic study on deformed and fractured stainless steels

with duplex structures was continued.

(3) The fatigue crack-growth rates and fracture toughness were measured

on the 25Mn-5Cr and 32Mn-7Cr austenitic steels.

(4) The elastic properties of the CF8M stainless-steel castings were measured ul trasonical ly.

(5) Mechanical property testing at 4 K in support of superconducting

magnet design and construction continued.

RESEARCH THIS YEAR (1983)

(1) The tensile properties and fracture toughness of 25Mn-5Cr weldments

produced by three welding processes will be measured at 4 K. (2) Fatigue crack-growth rates will be measured in CF8M castings as a function of ferrite content.

(3) Tensile tests will be conducted on ful 1-thickness stainless-steel

welds with the weld reinforcement in place.

(4) An evaluation of fully austenitic weld metals with low sulfur and

phosphorus contents will be initiated.

(5) A series of stainless steel welds with varying oxygen contents will

be produced.

(6) Mechanical property testing at 4 K in support of superconducting

magnet design and construction will continue.

197
mmsm

mm
CORRELATIONS OF FATIGUE CRACK GROWTH RATE PARAMETERS

AT CRYOGENIC TEMPERATURES*

Y.-W. Cheng and R. L. Tobler

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Fatigue crack growth rate data of ferritic steels, AISI 300 series

stainless steels, and austenitic steel welds at 295, 76, and 4 K were collected and analyzed in terms of the exponent in the Paris equation. The

data considered pertain to constant amplitude fatigue loading and stress ratio

equal to 0.1. It is observed that there is a linear relation between log C

and m, the parameters of the Paris equation. The exponent, m, is insensitive

to the variation of fracture toughness and yield strength of the material

except in low fracture toughness materials. Temperature effects on m are observed for ferritic steels but not for austenitic steels.

* To be published in Proceedings of the ICF International Symposium on Fracture Mechanics, November 22-25, 1983, Beijing, China.

199

The fracture mechanics approach to the evaluation of fatigue crack

growth behavior provides a rationale for material selection as well as design criteria against fatigue failure. For this reason, the materials

industry includes fatigue crack growth studies in the materials design and

characterization process for a wide variety of component--e.g. , aircraft

components, nuclear hardware, marine structures, piping, and bridges.

Over the past few years, the fatigue crack growth studies at the National

Bureau of Standards (NBS) have been used to characterize a wide variety of

cryogenic structural alloys and their welds. This work has been spurred

by developments in energy- related fields and by magnetic fusion energy

programs. This paper summarizes the observed correlations of fatigue

crack growth rate parameters of selected materials at room and cryogenic

temperatures in terms of the exponent, m, in the Paris equation,

da/dN = C( Ak) m

, where da/dN is the fatigue crack growth rate in mm/cycle,

AK is the stress intensity range in MPa/m, C and m are material constants.

MATERIALS

The materials studied include ferritic steels [1-4], wrought AISI 300

series stainless steels [5-12], and austenitic steel welds with various

welding processes and filler metals [12-17]. The AISI 300 series stainless

steels are commonly used Fe-Cr-Ni alloys. The austenitic steel weld data

include data on welds of the AISI 300 series stainless steels, as well as

Contribution of NBS; not subject to copyright in the U.S.

201
Table 1. Summary of regression statistics for data of Fig. 1.

Ferritic Steels AISI 300 Series

Stainless Steels

Austenitic Steel

Wei ds

slope -1.468 -1.527 -1.474

a* 0.109 0.0753 0.0646

intercept -3.749 -3.884 -4.247

0* 0.507 0.272 0.291

R** 0.966 0.946 0.970

**R: correlation coefficient; R = 1 implies a perfect correlation

additional Fe-Cr-Ni-Mn, Fe-Cr-Mn, and Fe-Cr-Ni alloys. The testing media

are laboratory air, liquid nitrogen, and liquid helium for test

temperatures of 295, 76, and 4 K, respectively. The data considered

pertain to constant amplitude fatigue loading with a stress ratio equal to

0 . 1 . The majority of the data collected were obtained in the laboratories

of NBS using compact specimens or bend specimens [2-4,6,7,9-11,13-17].

Additional relevant data [1,5,8,12] were included in the analysis for

comparison and confirmation. For details of specimen preparation,

material compositions and treatment, it is necessary to refer to the

original publications.

RESULTS AND DISCUSSION

The Relation Between C and m

It has been shown that there is an empirical linear relation between

log C and m for structural steels at room temperature [18,19]. This study

extended the correlation to ferritic steels, AISI 300 series stainless

steels, and austenitic steel welds at cryogenic temperatures . The results

are shown in Fig. 1. Least-squares regression analyses have been

performed for each of the three sets of data of Fig. 1, and the regression

statistics are summarized in Table 1. The results indicate that there is

a definite linear relation between log C and m for the materials studied.

A linear relation between log C and m has the following implications.

(1) The two constants (C and m) in the Paris equation are not independent

202
and reduce to one constant (C or m) for a specific class of material.

(2) Materials with higher values of m have a better fatigue crack growth

resistance than those with lower m values in the lower half of the stress

intensity range where the Paris equation applies, or vice versa at the

higher stress intensity range.

However, it should be pointed out that different materials with same

m values and different C values have been observed [20] and the empirical

relation between log C and m should be treated as a trend rather than a rule.

The Relation Between m and Fracture Toughness and Yield Strength

The results of m vs. fracture toughness and m vs. yield strength are plotted in Fig. 2. Figure 2a shows that the value of m is independent

of fracture toughness at fracture toughnesses greater than 90 MPa/m.

Below 90 MPa/m, high values of m are observed. The high values of m in

low fracture toughness specimens can be attributed to the concurrent

occurrence of monotonic fracture, such as cleavage or intergranular

fracture, with fatigue crack growth [20].

As shown in Fig. 2b, the value of m is independent of yield strength,

except that in high yield strength specimens the scatter in m becomes more pronounced.

The Relation Between m and Material Type

Examination of Fig. 1 has shown that there is a temperature effect on m for ferritic steels when temperature decreases from 295 to 76 and 4 K,

but not for AISI 300 series stainless steels and austenitic steel welds.

The effect of temperature in ferritic steels relates to the effect of

fracture toughness on m, as mentioned in the previous section. Ferritic

steels are b.c.c. metals that exhibit ductile-to-brittle transitions when

temperature decreases sufficiently. At low temperatures, ferritic steels

become brittle (low fracture toughness), and this produces high m values.

No temperature effect on m is observed in AISI 300 series stainless steels

and austenitic steel welds because these f.c.c. metals do not exhibit

ductile-to-brittle transitions as temperature is decreased from 295 to

4 K.

203
Fig. 1 also indicates that the AISI 300 series stainless steels have

least spread in m values and the austenitic steel welds have a higher

average m value than the AISI 300 series stainless steels do (4,43 vs. 3,57).

CONCLUSION

The following conclusions can be drawn from the present study:

1. There is a linear relation between log C and m in the Paris equation

for the cryogenic structural alloys and welds studied.

2. The value of m is independent of fracture toughness and yield

strength of the materials except in low fracture toughness materials;

in the latter case, m increases as fracture toughness decreases

below 90 MPa/m.

3. There is a temperature effect on m for ferritic steels but not for

austenitic steels; low temperatures tend to yield higher m values for

ferritic steels.

ACKNOWLEDGMENTS

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

Magnetic Fusion Energy.

REFERENCES

1. V. V. Pokrovski, Strength Mater. , 10 (5), 534 (1978).

2. R. L. Tobler, R. P. Mikesell, R. L. Durcholz, and R. P. Reed, in

Properties of Materials for Liquefied Natural Gas Tankage , ASTM STP

579, American Society for Testing and Materials, Philadelphia, PA

261 (1975).

3. R. L. Tobler, R. P. Reed, and R. E. Schramm, J. Eng. Mater. Techno! ., 100 , 189 (1978).

4. R. L. Tobler, R. P. Mikesell, and R. P. Reed, in Fracture Mechanics

(Eleventh Conference ), ASTM STP 677, American Society for Testing

and Materials, Philadelphia, PA, 85 (1979).

5. M. W. Mahoney and N. E. Paton, Nucl . Technol . , 23 53 (1974).

204
6. D. T. Read and R. P. Reed, in Materials Studies for Magnetic Fusion

Energy Applications at Low TemperaturesI

, NBSIR 78-884, National

Bureau of Standards, Boulder, Colorado, 93 (1978).

7. D. T. Read and R. P. Reed, in Metal Science of Stainless Steels , Metallurgical Society of New York, 92 (1979).

8. F. R. Schwartzberg, in Materials Research for Superconducting

Machinery- 1 1 , Semi-annual Technical Report, ARPA-2569 , National

Bureau of Standards-Advanced Research Projects Agency (1974).

Available from NTIS, Springfield, Virginia.

9. R. L. Tobler and R. P. Reed, in Advances in Cryogenic Engineering , 22, 35 (1977).

10. R. L. Tobler and R. P. Reed, in Advances in Cryogenic Engineering , 24, 82 (1978).

11. R. L. Tobler and R. P. Reed, in ASTM 15th Natl. Symp. on Fracture

Mechanics , College Park, Maryland (1982).

12. J. M. Wells, R. Kossowsky, W. A. Logsdon, and M. R. Daniel, in

Materials Research for Superconducting Machinery- IV, Semi-annual

Technical Report , ADA036919, National Bureau of Standards-Advanced

Research Projects Agency (1976). Available from NTIS, Springfield,

Virginia.

13. H. I. McHenry, J. W. Elmer, and T. Inoue, in Materials Studies for

Magnetic Energy Applications at Low TemperaturesV

, NBSIR 82-1667,

National Bureau of Standards, Boulder, Colorado, 245 (1982).

14. H. I. McHenry and T. A. Whipple, in Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures- 1 II , NBSIR 80-1627,

National Bureau of Standards, Boulder, Colorado, 155 (1980).

15. R. L. Tobler, H. I. McHenry and R. P. Reed, in Advances in Cryogenic

Engineering , 24 , 560 (1978).

16. T. A. Whipple and H. I. McHenry, in Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures- IV , NBSIR 81-1645,

National Bureau of Standards, Boulder, Colorado, 273 (1981).

17. T. A. Whipple, H. I. McHenry, and R. T. Read, Weld. J. , 60 (4), 72s

(1981).

18. K. Tanaka and S. Matsuoka, Int. J. Fract. , 13 563 (1977).

19. T. R. Gurney, Fatigue of Welded Structures , 2nd edition, Cambridge

University Press (1979).

20. R. L. Tobler and Y. W. Cheng, to be presented at the Symposium on Fatigue at Low Temperatures , Louisville, Kentucky (May 1983).

205
List of Figures

Figure 1(a). Relation between log C and m for the AISI 300 series stainless

steel s

Figure 1(b).

Figure 1(c).

Figure 2(a).

Figure 2(b).

Relation between

Relation between

log C and m for

log C and m for

austenitic steel

ferritic steels.

welds.

Relation between m and fracture toughness. Note: fracture

toughnesses were converted from the J- integral which was obtained from the single specimen test technique.

Relation between m and yield strength.

206
-5

-7

O

o>o

AISI 300 Series Stainless Steels

o 295 K

• 76K

a 4 K

2 4 6 8

m

Figure 1(a). Relation between log C and m for the AISI 300 series stainless

steels.

207
-5

-7

-9

O

o»o

-13

-15

Austenitic Steel Welds

o 295 K

• 76K

a 4 K

_L

2 4 6 8

m

Figure 1(b). Relation between log C and m for austenitic steel welds.

208
Figure 1(c). Relation between log C and m for ferritic steels.

209
o *

CO a.

co

co

LUzXCD3oH

LU

CCO<

GC

LL

UJ

210

Figure

2(a).

Relation

between m

and

fracture

toughness.

Note:

fracture

toughnesses

were

converted

from

the

J-integral

which

was

obtained

from

the

single

specimen

test

technique.
oocDoo

CMoo

00ooO

Ul

-Mcnc

CL) i-

+->

(/)

C\J

<u

s13CD

Ll_

211

YIELD

STRENGTH,

MPa

DEFORMATION AND FRACTURE OF STAINLESS STEEL

CASTINGS AND WELDMENTS AT 4 K*

T. A. Whipple and E. L. Brown

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

A number of stainless steel weldments and castings that were deformed at

4 K have been examined with optical and electron microscopy. The purpose of

this investigation was to assess the effects of residual delta-ferrite on the

deformation and fracture mechanisms of these materials at cryogenic temperatures.

Weldments and castings have very similar microstructures, but the deltaferrite in castings is much coarser as a result of slower solidification and

cooling rates. This wide variation in the scale of the structure provides a good opportunity for the study of structure and property relationships in

these materials.

It has been found that deformation and fracture are greatly influenced by

the ferrite morphology. The study of the coarser structure in castings has

allowed more direct observation of fracture, which has led to some preliminary

conclusions concerning fracture mechanisms. Cleavage fractures of deltaferrite particles have been observed. It appears that these fractures,

although not forming a continuous crack path, assist in the void formation in

the ductile austenite matrix. Another observed crack propagation mechanism is

that of voild nucleation near the austenite delta-ferrite interface. Although

interface separation has not been observed, the inhomogeneous deformation near

the interface appears to cause premature void nucleation and growth. The

occurrence of one of the above mechanisms rather than the other depends on a number of factors, including ferrite morphology and the properties of the two

phases. The morphology and properties, in turn, depend on both the alloy

composition and solidification conditions.

* Published in "Trends in Welding Research in the U.S.," American Society for

Metals Conference Proceedings, New Orleans, LA, November 16-18, 1981,

Editor, S. A. David (American Society for Metals, Metals Park, OH, 1982).

213

INTRODUCTION

The structural components for most superconducting

magnets are being made of stainless steels. Because of the

large structures, high stresses and low operating temperature,

welded construction is the most attractive fabrication

method. Since the structures are being made of thick plate,

fracture toughness is an important consideration. It has

been found that stainless steel weldments generally have a much lower cryogenic fracture toughness than the stainless

steel base metal (1,2).

To date, three factors have been established that have

an effect on the cryogenic fracture toughness of stainless

steel weldments: the delta-ferrite content of the weldment,

the nitrogen concentration in the weld, and the fusion zone grain size of fully austenitic weldments (1,2, 3, 4). Nevertheless, it still is not possible to predict the toughness of a stainless steel weldment on the basis of these factors.

Apparently there are factors beyond those stated above that cause the observed wide variations in fracture toughness

(2,5,6). This paper contains the results of an investigation

that is intended to define more precisely the role of deltaferrite and to identify other factors that have an influence

on the cryogenic toughness of stainless steel weldments and

castings.

In this investigation qualitative observations were made

on various stainless steel weldments and castings which have

been tensile and fracture toughness tested at 4K. Also, a series of CF8M castings, with varying delta-ferrite contents,

have been tested at 4K. There are presently very few 4 K

mechanical property data on castings, but a few initial tests

have shown a wide variation in fracture toughness similar to

weldments (5). This is not surprising because weldments are actually rapidly solidified castings. Stainless steel

castings also generally contain a certain percentage of

delta-ferrite, although on a much coarser scale than that of

weldments, owing to the slower solidification of castings.

Because of the coarser structure and less directional nature

of the delta-ferrite in castings, they provide an excellent

model material for the study of basic deformation and fracture

mechanisms in duplex austenite/delta-ferrite materials. It

is anticipated that this study will aid in understanding the causes of the extreme property variations in weldments.

The first part of this investigation was a metallographic and fractographic investigation of stainless steel

215
weldments and castings, containing delta-ferrite, which were deformed and fractured at 4 K. Secondly, mechanical property

tests were performed on a controlled series of CF8M stainless

steel castings with varying ferrite contents.

MATERIALS

The materials for the initial part of this study are stainless steel weldments and castings that have been tensile

and fracture toughness tested at 4 K. These materials did

not come from a controlled series, however, they all fall

into the AISI type 304 or 316 stainless steel composition

ranges. These variations in composition, as well as variation

in the casting or welding process, are appropriate since the

intent of this study was to investigate the possible deformation and fracture mechanisms in duplex austenite delta-ferrite

structures at 4 K and not just the mechanism for a specific

material. The welding or casting process, delta-ferrite

content and alloy are noted within the text as needed.

The chemical composition and delta-ferrite content of

the controlled series of CF8M castings are presented in Table

1. The castings were produced by a commercial vendor, and

the procedures used have been reported previously (7). The

castings were received in the form of 30.5 cm (12 in) X 61.0 cm

(24 in). X 10.2 cm (4 in) blocks. All specimens were

removed from areas at least 15.2 cm (6 in) from either end

and 3.8 cm (1.5 in) from either surface. This was done to

avoid any effects of the columnar grain structure near the

casting surfaces.

Table 1. Chemical analysis of five heats of CF8M with variable ferrite content.

Heat C Mn Si Cr Ni Mo N % Delta-Ferrite*

1 0.06 1.10 0.59 18.04 13.24 2.12 .05 1.1

2 0.06 0.38 1.16 19.62 10.61 2.12 .05 8.2

3 0.06 0.36 1.13 19.23 8.25 2.16 .05 17.2

4 0.06 0.35 0.95 20.40 8.22 2.13 .05 23.1

5 0.05 0.39 1.11 22.54 9.84 2.26 .05 28.5

‘Measured by a metal lographlc point count.

216
PROCEDURES

Metal lographic Specimen Preparation

For this investigation a number of sectioning, metallographic and fractographic techniques were employed. Each

technique allows the observation of different phenomena.

Macroscopic examination, optical microscopy, and scanning

electron microscopy (SEM) were used for making observations.

For the most part metallographic specimens were prepared

by using standard mechanical grinding and polishing techniques, but in a few cases electrolytic polishing was used.

A 10% oxalic acid electrolytic etch was used as a general

purpose etch to reveal delta-ferrite morphology and grain

boundaries. A 10 N KOH electrolytic etch was used to reveal

the delta-ferrite structure only. This is a depostion etch

that colors the delta-ferrite light blue to brown. The main

attribute of this etchant is that it does not remove any

material; therefore, it will not round the edges of cracks or produce pitting and other etching artifacts. A boiling

solution of 1 part H

? 0? and 2 parts 1LP0. was used to reveal

strain induced martensite.

Tensile tests were conducted at 4 K using 2.5-cm (1-in)

gage length and 0.64-cm (0.25-in) diameter specimens. The

test procedures and the cryostat for the low temperature

tests have been described by Reed (8).

Fracture toughness tests were conducted at 4 K using

three-point bend specimens with a 2.54-cm (1-in) width and

2.54-cm (1-in) thickness. The single-specimen J-integral

procedure (9) was used to determine J, , the value of J at

the onset of crack extension. The plane strain fracture

toughness, K. (J), was determined from J, using the relationship proposed c by Begley and Landes (10):

lc

where E is Young's modulus and v is Poisson's ratio. Tests

were conducted in an electrohydraul ic fatigue machine

Mechanical Testing

217
equipped with the cryostat described by Fowlkes and Tobler

(11). The test data is collected and reduced by a laboratory

minicomputer. This system has been described by Read (12).

RESULTS AND DISCUSSION

Low Temperature Deformation

It is generally true that bcc materials, such as deltaferrite, go through a ducti le-to-brittle transition at low

temperatures, whereas fee materials, such as austenite, do

not . The temperature at which this transition occurs in

steels is influenced mainly by alloy content and microstructure; a higher nickel content and a smaller grain size lead

to a lower transition temperature. Another factor to consider

is the increase in strength with a reduction in temperature;

bcc materials have a much more dramatic increase in yield

strength than fee materials at low temperatures.

These general observations lead to some expectations for

the behavior of duplex stainless steels. At higher temperatures, both austenite and ferrite are ductile and have

similar strengths. In this case the amount and morphology of

delta-ferrite is expected to have a much smaller effect on strength and ductility than at low temperatures, when the

del ta-ferrit.e is strong and brittle. The effect of the

property transition in the delta-ferrite should be to increase

the strain hardening rate. Data that demonstrate the greater

strengthening effect of delta-ferrite at cryogenic temperatures

have been presented by Read et al.(l).

Figures 1 and 2 are micrographs of 316L stainless steel

weld metal tensile specimens that were tested at room temperature and 4 K, respectively. Note that at room temperature

(Figure 1) the delta-ferrite is highly elongated in the

direction of the tensile axis and at 4 K (Figure 2) the

ferrite maintains its original morphology and shows no indication of elongating with the matrix.

During deformation at low temperatures, austenite can go

through a strain induced martensitic transformation. The

degree of transformation and the critical strain at which it occurs are a function of the alloy content. Although the

stability of the austenite matrix will quantitatively affect

the strength and toughness of duplex austenitic stainless

steel castings and weldments, it is only of interest in this

qualitative discussion in that it provides a method of

observing deformation in the austenite.

218
Figure 1. Micrograph of a 316L submerged-arc weld metal tensile specimen, tested at room temperature. KOH etch.

Figure 2. Micrograph of a 316L submerged-arc weld metal tensile specimen, tested at 4K. KOH etch.

219
Figures 3 and 4 are micrographs of areas near the crack

tips in two CF8* stainless steel castings tested at 4 K.

The intent of these micrographs is to show the inhomogeneity

in deformation caused by the presence of delta-ferrite.

Figure 3 is from a 0 % del ta-ferrite -casting; the straight

lines near the crack tip are strain-induced martensite. Note

that the lines are straight and fairly uniform, only changing

direction at grain boundaries. This can be contrasted with

Figure 4, which is a similar region in a 14.5% delta-ferrite

casting. In this case, the martensite formation is nonuniform

and seems to be concentrated in association with the deltaferrite morphology. Another example of the effect of the

presence of delta-ferrite on deformation patterns in the

austenite matrix is shown in Figure 5. In Figure 5, bending

of martensite bands near the delta-ferrite particles can be seen in a CF8M* casting with 9 % ferrite, indicating that the

delta-ferrite is nondeforming and the matrix is flowing

around it.

Evidence to this point has indicated that delta-ferrite

does not deform at 4 K. However, work on the series of CF8M

castings has shown this to not be strictly true. As can be seen in Figure 6, the delta-ferrite does deform by a twinning

mechanism. The jogs in the delta-ferrite/ austenite interface and the fact that the twins are not observed in un- deformed structures demonstrate that they are in fact deformation twins. Further evidence for this conclusion is given in

Figure 7 in which three twin variants are shown. A stereographic projection of the angles between the three variants

shows that they are consistent with (112) twin planes, which

is typical for deformation twinning in ferritic materials.

In a macroscopic sense the deformation associated with

twinning in the delta-ferrite is small; however it may have

some significant secondary effects. One of these effects

can be seen in Figure 8. This figure shows bands of intense

martensite formation in the austenite at the ends of deformation twins in the delta-ferrite. This intense deformation

could create preferred sites for void nucleation in the

austenite.

The results of tensile tests on the five CF8M castings

at 4 K, are given in Table 2 and shown graphically in

*CF8 and CF8M are ASTM A351 grades of cast stainless steel

and are analogous to AISI 304 and 316, respectively.

220
Figure 3. Area near the crack tip of a 0% delta-ferrite CF8

casting. Oxalic acid etch.

Figure 4. Area near the crack tip of a 14.5% delta-ferrite

CF8 casting. KOH and oxalic acid etch.

221
Figure 5. Micrograph of a 9% delta-ferrite CF8M tensile specimen, deformed at 4K. H^PO^ etch.

Figure 6. Micrograph of the 28.5% delta-ferrite CF8M casting,

showing the deformation twinning that occurs at 4K. KOH etch.

222
Figure 7. An 8.2% delta-ferrite CF8M casting, showing three

deformation twin variants. H^PO^ etch.

Figure 8. Same specimen as Figure 7, showing austenite deformation associated with twinning in the del ta-ferrite. H„P0,

etch.

’3 u

4

223
Table 2. Mechanical properties of CF8M castings at 4K.

Heat Yield Strength

MPa U.T.S.

MPa

Elong.

%

K Ic

(J:

1 581 1215 48 MPa/m

352

2 653 986 21 148

3 641 1143 19 113

4 707 1221 19 106

5 851 1164 11 105

PLASTIC STRAIN, %

Figure 9. Stress versus plastic strain curves for the five

CF8M castings, at 4K.

224
Figure 9, as stress versus plastic strain. The important

feature to note on the stress-strain curves is the significant increase in strain hardening rate at low strains with

increasing delta-ferrite content. As mentioned previously

this is due to the increasing density of relatively non- deforming delta-ferrite particles which act as barriers to

dislocation motion.

Low Temperature Fracture

As a result of the above discussion on low-temperature

deformation, it appears appropriate to classify duplex

austenitic stainless steel, at low temperatures, as a ductile

matrix composite with hard, brittle discontinuous fibers.

This is not an exact classification, since the delta-ferrite

does show some evidence of deformation and can be either

continuous or discontinuous depending on solidification and

cooling conditions, but this classification does provide a good starting point for the discussion of fracture mechanisms.

There are three possible mechanisms for the fracture of

ductile matrix/hard-discontinuous fiber composites (13).

1. Brittle fracture of the fibers.

2. Matrix/fiber interface separation.

3. Fiber pull-out.

Brittle fracture of the fibers occurs when the local

stress in the fiber exceeds its fracture strength. The local

stress in the fiber is not necessarily the same as the

macroscopic stress on the sample; it can be considerably

higher owing to inhomogeneous plastic deformation in the

matrix. Interface separation is an important consideration

with artificial composites and with metal -matrix/particulate

composites; however, the interface between austenite and

ferrite appears to be very strong and no evidence of interfacial separation has been observed. Fiber pull-out is

sometimes associated with interface separation, but for the

purpose of this discussion it means a ductile failure of the

matrix near the matrix/fiber interface. This occurs when

the fiber does not fracture, but the highly inhomogeneous

deformation of the matrix near the fiber results in void

nucleation at this location. The occurrence of one of these

mechanisms over the other depends on fibre strength, content

and morphology.

In this study evidence was found for both fiber fracture

and fiber pull-out. Until this investigation there had been

225
no direct observation of the brittle fracture of deltaferrite in weldments and castings at 4 K. Although this had

long been assumed to be the cause of the reduction in toughness

of stainless steel weldments with increasing delta-ferrite

content, direct observation was difficult because of the

extremely small size of the delta-ferrite in weldments. In

this investigation observation were made on both castings and

weldments made with high heat input processes. The much

slower solidification and cooling rates of these materials

results in a coarser structure, which is more easily observed.

With these materials extensive evidence was found of

brittle fracture of delta-ferrite in weldments and castings

deformed at 4 K. Evidence for cleavage fracture of deltaferrite is shown in Figure 10 (a micrograph from the 14.5%

delta-ferrite CF8 casting). This micrograph shows consistent

angular relationships between the fractures, indicating that

fracture is occurring on a specific set of crystallographic

planes.

Weldments differ from castings mainly in that the

solidification and cooling rates are much greater. This

results in a more directional solidification pattern and a much finer structure. The greater cooling rate results in

less diffusion during the del ta-ferri te-to-austeni te transformation causing the delta-ferrite to be more closely associated

with the original dendritic solidification pattern. When

fracture is associated with cracking of the residual deltaferrite, having continuous delta-ferrite can provide a preferred crack path and lower toughness than occurs for a material of the same delta-ferrite content in a discontinuous

morphology.

An example of this preferred crack path is shown in

Figure 11 in the crack profile of a fracture toughness

specimen from a 316L electroslag weldment. An electroslag

weldment experiences a cooling rate greater than that of a casting, but much less that that of other welding processes.

This results in a relatively coarse but continuous

delta-ferrite morphology. Figure 11 shows that the

dendritic growth direction determined the crack propagation

direction even though the stress state was such that the

crack should propagate straight. The crack was propagated

down dendrite cores and secondary dendrite arms.

Further evidence that the delta-ferrite in this

weldment fractured in a brittle manner is provided by the

SEM fractograph in Figure 12. This fractograph also

226
Figure 10. Area near the crack of a 14.5% delta-ferrite CF8

casting. KOH and oxalic acid etch.

Figure 11. Crack profile of a 316L electroslag weldment.

Oxalic acid etch.

227
demonstrates a mixture of brittle delta-ferrite fracture and

ductile austenite failure, which will be discussed later.

As mentioned previously, two possible discontinuous

fiber composite fracture mechanisms are of interest: fiber

fracture and fiber pull-out. The evidence already presented

has shown that fiber fracture does occur in coarse duplex

structures. In the case of weldments with much finer structures, limited evidence has been found for both fracture

mechanisms. Figure 13 is a fractograph of the fracture

surface of a 308L FCMA weldment. The area indicated by the

arrow appears to be a subsurface delta-ferrite dendrite,

where the inhomogenous deformation nucleated a ductile

failure near the delta-ferrite/austenite interface. Composite fracture theory predicts that, for a discontinuous

fiber composite, there is a critical aspect ratio (fiber

length-to-diameter ratio) at which a transition from fiber

fracture to fiber pull-out fracture mechanisms occurs; a high

aspect ratio favors fiber fracture. The aspect ratio of

fiber is dependent on orientation to the tensile axis: the same fiber will have a high aspect ratio if it is aligned

with the tensile axis and a low one if it is perpendicular to

the tensile axis. In Figure 13 it is apparent that the

dendrite was perpendicular to the tensile axis; therefore

fiber pull-out occurred.

From the observations made to date fiber pull-out

appears to be the less prevalent of the two fracture mechanisms observed. The most commonly observed mechanism is

brittle fracture of the delta-ferrite. If the delta-ferrite

is continuous the brittle fracture will result in specimen

failure. However, in the case of discontinuous delta-ferrite

the brittle fracture of discrete particles will cause stress

concentrations that will result in ductile failure of the

austenite matrix. An example of this mechanism is shown in

Figure 14. This micrograph shows brittle fractures in a delta-ferrite particle with ductile fracture in the austenite

between two cracks. This mechanism will result in a fracture

surface displaying a mixed mode fracture. This type of

fracture has been shown previously (Figure 12) and is commonly

observed in these materials.

The deformation twinning in the delta-ferrite does have

an effect on the fracture process. The different orientation

of the cleavage planes in the twins causes the crack to

change direction. This can be seen in the micrograph in

Figure 15. The deformation twinning and the change in crack

path will result in additional energy being absorbed in the

228
Figure 12. SEM fractograph of the fracture surface of a 316L electroslag weldment.

Figure 13. SEM fractograph of a 308L flux - core - metal - arc weldment229
Figure 14. Micrograph of the 28.5% delta-ferrite CF8M casting, showing the combination of ductile and brittle failure.

H^PO^ etch.

Figure 15. Same specimen as Figure 14, showing the interaction

of the cracks with deformation twins. H^PO^ etch.

230
fracture process. This will result in a higher fracture

toughness. It is not possible, at this time, to determine

the quantitative effect that this has on the fracture toughness

of these materials.

The results of the fracture toughness testing, at 4 K, on the five CF8M castings are shown in Figure 16. This shows

a decrease in fracture toughness with increasing delta-ferrite

content up to about 15% delta-ferrite. At higher deltaferrite contents the fracture toughness is essentially

constant. Metallographic and fractographic examination has

shown this to be related to the establishment of a continuous

delta-ferrite crack path at this delta-ferrite content.

Castings with less than 15% deltaferrite show a mixed-mode

ductile/brittle fracture, and those above 15% delta-ferrite

show a completely brittle fracture.

Figure 16. The effect of delta-ferrite on the fracture toughness of CF8M castings and 316L weldments.

231
SUMMARY

Samples of stainless steel castings and weldments that

had been deformed and fractured at 4 K were examined. The

thrust of the investigation was to determine the mechanisms

by which residual delta-ferrite affects the deformation and

fracture of duplex austenite/delta-ferrite structures at

cryogenic temperatures. Evidence was presented that indicates that at low temperatures the delta-ferrite acts as a hard, brittle second phase particle. The manner in which

these particles affect the mechanical properties of the

material depends on the amount, properties and morphology of

the second-phase particles. It was observed that the deltaferrite reinforces the austenite matrix and causes the

deformation to be inhomogeneous. Evidence was found for two

fracture mechanisms: (1) brittle fracture of the deltaferrite mixed with ductile failure of the matrix between

delta-ferrite fractures and (2) ductile failure in the matrix

near the austenite/delta-ferrite interface, which was caused

by the highly inhomogeneous deformation in this region. The

occurrence of one of these mechanisms over the other depends- on the morphology and properties of the delta-ferrite. It was also observed that delta-ferrite deforms by a twinning

mechanism at 4 K. These deformation twins have an effect on the crack propagation.

REFERENCES

(1) D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D.

Thomas, "Metallurgical Factors Affecting the Toughness

of 316L SMA Weldments at Cryogenic Temperatures," Weld.

J. Res. Suppl

, 59, 1980, p. 104s.

(2) T. A. Whipple, H. I. McHenry, and D. T. Read,

"Fracture Behavior of Ferrite-Free Stainless Steel

Welds in Liquid Helium," Weld. J. Res. Suppl., 60,

(4), 1981, P. 72s.

(3) E. R. Szumachowski and H. F. Reid, "Cryogenic

Toughness of SMA Austenitic Stainless Steel Weld

Metals: Part I - Role of Ferrite," Weld. J. Res.

Suppl

, 57, 1978, p. 325s.

(4) E. R. Szumachowski and H. F. Reid, "Cryogenic

Toughness of SMA Austenitic Stainless Steel Weld

Metals: Part II - Role of Nitrogen," Weld. J.

Res. Suppl. , 58, 1979, p. 34s

232
(5) T. A. Whipple and H. I. McHenry, "Evaluation of

Weldments and Castings for Liquid Helium Service,"

in Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures-IV, NBSIR 81-1645,

1981, p. 273.

(6) T. A. Whipple and D. J. Kotecki , "Weld Process

Study for 316L Stainless Steel Weld Metal for

Liquid Helium Service," in Materials Studies for

Magnetic Fusion Energy Application at Low

Temperatures-IV , NBSIR 81-1645, 1981, p. 303.

(7) L. Finch, " The Influence of Ferrite and Nitrogen

on Mechanical Properties of CF8M Cast Stainless

Steel," In Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures-IV, NBSIR

81-1645, 1981, p. 337.

(8) R. P. Reed, "A Cryostat for Tensile Tests in the

Temperature Range 300 to 4 K," in Advances in

Cryogenic Engineering, VoT. 7 , Plenum Press, New

York, 1962, pp. 448-454.

(9) G. A. Clarke, W. R. Andrews, P. C. Paris, and D.

W. Schmidt, "Single Specimen Tests for J, Determination," in Mechanics of Crack Growth , lc Proceedings of the 1974 National Symposium on Fracture

Mechanics, ASTM STP 590, American Society for

Testing and Materials, Philadelphia, PA, 1976,

pp. 27-42.

(10) J. A. Begley and J. D. Landes, "The J-Integral as

a Fracture Criterion", in Fracture Toughness , Proceedings of the 1971 National Symposium on Fracture Mechanics, Part II, ASTM STP 514 American

Society for Testing and Materials, Philadelphia,

PA, 1972, pp. 1-20.

(11) C. W. Fowlkes and R. L. Tobler, "Fracture Testing

and Results for a Ti 6A1-4V Alloy at Liquid Helium

Temperatures," Eng. Fract. Mech. , 8, 1976,

pp. 487-500.

(12) D. T. Read, "The Computer-Aided J-Integral Test

Facility at MBS," in Materials Studies for

Magnetic Fusion Energy Applications at Low

Temperatures-III , NBSIR 80-1627, 1980, p. 205.

233

MECHANICAL PROPERTIES OF STAINLESS STEEL CASTINGS AT 4 K*

T. A. Whipple and H. I. McHenry

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

In some cryogenic structural applications it would be advantageous to use stainless steel castings, rather than wrought and welded structures. However,

to date there are very few data on the mechanical properties of stainless

steel castings at cryogenic temperatures. Stainless steel castings usually

contain a certain amount of delta-ferrite to prevent hot cracking. The

present program was designed to evaluate the effects of delta-ferrite and

interstitial nitrogen on the strength and fracture toughness of stainless

steel castings at 4 K.

Ten heats of cast CF8M stainless -steel , which is the casting equivalent

of AISI 316, were produced by a commercial vendor. In five of the heats the

chromium and nickel contents were varied to obtain different delta-ferrite

contents, ranging from 1.1 to 28.5 percent. The other five heats all had

approximately 8 percent delta-ferrite and the nitrogen content was varied from

0.02 to 0.20 percent. As was expected from previous data on weldments and

wrought materials, an increase in either delta-ferrite or nitrogen content

increases the -yield strength. In the five heats with varying nitrogen content

the increase in strength resulted in a corresponding decrease in fracture

toughness. An increase in del ta-ferrite content also gave a decrease in

fracture toughness up to approximately 15 percent delta-ferrite. Above this

level the fracture toughness remained constant with increasing del ta-ferrite

content. This has been shown, by metal lographic and fractographic evidence,

to be due to the establishment of a continuous delta-ferrite crack path, at

approximately 15 percent del ta-ferrite.

+ Work supported by Office of Fusion Energy, Department of Energy

* To be published in Austenitic Steels at Low Temperatures , Eds. R. P. Reed

and T. Horiuchi, Plenum Press, New York (1983).

235

INTRODUCTION

The limited use of stainless steel castings for 4-K service is

partially due to the limited data available on their mechanical properties at 4 K. 1 The most significant 4-K data available on the

tensile and fracture toughness properties of stainless steel

castings came from a series of CF8 centrifugal castings with varying

delta-ferrite contents. 1

* 2 There were six heats from three different vendors with delta-ferrite contents ranging from 0 to 14.5

percent. The results generally showed an increase in strength and a decrease in fracture toughness with increasing ferrite content.

The present study was undertaken to evaluate the effects of

delta-ferrite and nitrogen on the 4-K tensile and fracture toughness

properties of a series of nine CF8M stainless steel castings.

MATERIALS

The materials used in this investigation were nine CF8M stainless steel castings whose chemical composition and delta-ferrite

content are given in Table 1. For the first five castings, the

nickel and chromium contents were varied to obtain five residual

delta-ferrite contents, ranging from 1.1 to 28.5 percent. For the

remaining castings, the nitrogen content was varied and the other

elements were balanced .to obtain a constant delta-ferrite content.

The castings were produced by a commercial vendor; the procedures used and room temperature mechanical properties have been

reported previously. 3 The castings were received in the form of

30.5-cm x 61.0-cm x 10.2-cm blocks. All specimens were removed from

237
Table 1. Chemical composition* and delta-ferrite

content of CF8M castings.

Vol

Percent

Heat No c Mn Si Cr Ni Mo N Deltai-Ferrite

47640 0.06 1 . 10 0.59 18.04 13.24 2.12 0.05 1.1 48988 0.06 0.38 1.16 19.62 10.61 2.12 0.05 8.2 47521 0.06 0.36 1.13 19.23 8.25 2.16 0.05 17.3 47552 0.06 0.35 0.95 20.40 8.22 2.13 0.05 23.1 49021 0.05 0.39 1.11 22.54 9.84 2.26 0.05 28.5

49020 0.07 0.39 1.09 19.32 10.00 2.07 0.02 8.0

47364 0.06 0.40 1.30 19.17 9.14 2.00 0.10 8.5 47390 0.07 0.41 0.98 21.00 9.46 2.13 0.144 9.5

47399 0.06 0.33 1.35 21.66 9.68 2.83 0.20 9.5

*Chemical composition in weight percent

balance Fe.

areas at least 15.2 cm from either end and 3.8 cm from either sur face. This was done to avoid any effects near the casting surface.

The test specimens were 0.635-cm smooth bar tensile specimens, with

a 2.54-cm gage length.

PROCEDURES

The tensile tests at 4 K were conducted with the test apparatus

and cryostat described by Reed. 4

For the fracture toughness tests, a test procedure using threepoint-bend specimens has been developed. 5 The single-edge-notched

three-point-bend specimens were 2.54-cm wide, 2.54-cm thick, with a span of 10.2 cm. A direct measurement of the load-line displacement

was made using the comparison bar method developed by Dawes. 5 The

test data were acquired and reduced by a laboratory minicomputer, as described by Read. 7

The volume-percent delta-ferrite was measured metallographically, using a computerized image analyzer. The specimens were electrolytically etched with a 10 N KOH solution, which gives good

phase contrast. The values reported for volume percent represent

the average for 400 fields at 100 times magnification.

RESULTS AND DISCUSSION

The results of the 4-K tensile tests are given in Table 2 and

shown graphically in Figures 1 and 2. Figure 1 contains the results

from the heats with a constant nitrogen content and variable deltaferrite contents. The yield strength increased with increasing

delta-ferrite content. The tensile strength did not show a definite

trend with respect to the delta-ferrite content, owing to the

decrease in ductility with increasing delta-ferrite.

At low temperature the delta-ferrite, which has a bcc crystal

structure, goes through a ductile-to-brittle transition. At 4 K, it

238
Table 2. Mechanical properties of CF8M castings at 4 K.

Heat No. Yield Strength

MPa Tensile Strength MPa Elongation

% Reductionin Area

K, ( J)

M^a/m

47640 581 1215 49 33 352*

48988 653 986 21 15 148 47521 641 1143 19 13 113 47552 707 1221 19 14 106 49021 851 1164 11 17 104 49020 650 1122 22 14 129 47314 725 1152 20 17 126 47390 951 1086 13 15 74 47399 1100 1138 6 4 94 *Invalid by specimen size requirements

.

is essentially nondeforming, whereas the fee austenite remains

ductile. Therefore, an increase in the delta-ferrite content causes

a dramatic increase in the strain-hardening rate of the duplex material. The yield strengths of the 8.2- and 17.3-percent deltaferrite castings were found to be the same. Further testing in the

low-strain region has shown this to be due to a difference in the

proportional limit of the austenite matrix. 5 The 17.3-percent

delta-ferrite casting had a lower proportional limit and a higher

strain-hardening rate.

Fig. 1. The effect of delta-ferrite on the tensile properties

of castings containing 0.05 percent nitrogen.

239
NITROGEN, wt. %

Fig. 2. The effect of nitrogen on the fracture toughness

of castings containing 9 ± 1 percent delta-ferrite.

Figure 3 shows the effect of nitrogen on the mechanical properties of CF8M stainless steel castings at 4 K. An increase in the

nitrogen content increased the yield strength and reduced the ductility. Again, the tensile strength did not show a trend owing to

the difference in ductility. The effect of nitrogen on the tensile

properties was small at nitrogen contents below 0.10 percent.

The results of the fracture toughness tests at 4 K are given in

Table 2 and shown graphically in Figures 3 and 4. The fracture

toughness result on the 1.1-percent delta-ferrite casting is invalid

by specimen size requirements and is used here only as an approximate value. However, it can be said that the fracture toughness of

this alloy is much higher than the other four heats and is comparable to that of wrought materials of similar compositions. The

interesting feature of the fracture toughness data is that the three

high delta-ferrite heats (17.3, 23.1, and 28.5 percent) have approximately the same fracture toughness. Metallographic and fractographic observations have shown that this is due to the establishment of continuous delta-ferrite crack path at high delta-ferrite

contents

Figure 4 shows the fracture toughness as a function of nitrogen

content. There is a general trend toward decreasing fracture tough- ness with increasing nitrogen content. The significant scatter in

240
these results is probably caused by variations in the delta-ferrite

content of the castings.

CONCLUSIONS

1. An increase in either the delta-ferrite or nitrogen content of

CF8M stainless steel castings increases the yield strength and

reduces the ductility at 4 K.

2. Increasing the delta-ferrite content of CF8M castings decreases

the fracture toughness at 4 K, up to a delta-ferrite content of

17 percent. Above this point, the fracture toughness is constant.

3. Increasing the nitrogen content of CF8M castings above 0.10 percent reduced the fracture toughness.

Fig. 3. The effect of delta-ferrite on the fracture toughness

of castings containing 0.05 percent nitrogen.

Fig. 4. The effect of nitrogen on the fracture toughness of

castings containing 9 ± 1 percent delta-ferrite.

241
ACKNOWLEDGMENTS

The authors thank L. Finch of ESCO Corporation for his work in

designing the alloy series and R. J. Fields of the National Bureau

of Standards for the quantitative metallography work he performed.

This research was supported by the United States Department of

Energy, Office of Fusion Energy.

REFERENCES

1. D. Dew-Highes and K. S. Lee, The choice of steel for the

Isabelle magnet tubes, in "Advances in Cryogenic

Engineering—Materials," Vol. 26, Plenum Press, New York

(1981), pp. 151-157.

2. E. L. Brown, T. A. Whipple, and R. L. Tobler, Fracture toughness

of CF8 castings in liquid helium, to be published in Metall . Trans .

3. L. Finch, The influence of ferrite and nitrogen on the mechaniproperties of CF8M cast stainless steel, in "Materials

Studies for Magnetic Fusion Energy Applications at Low Temperatures—IV," NBSIR 81-1645, National Bureau of Standards,

Boulder, Colorado (1981), pp. 337-355.

4. R. P. Reed, A cryostat for tensile tests in the temperature

range 300 to 4 K, in "Advances in Cryogenic Engineering,"

Vol. 7, Plenum Press, New York (1961), pp. 448-454.

5. T. A. Whipple, The effect of delta-ferrite on the mechanical

properties of CF8M stainless steel castings at 4 K, in

"Materials Studies for Magnetic Fusion Energy Applications at

Low Temperatures—V," to be published by the National Bureau

of Standards, Boulder, Colorado.

6. M. G. Dawes, Welding Institute Research Bulletin No. 7 (1976),

p. 185.

7. D. T. Read, The computer-aided J-integral test facility at NBS,

in "Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures—III," NBSIR 80-1127, National Bureau of

Standards, Boulder, Colorado (1980), p. 205.

242
METALLOGRAPHY OF DUPLEX STAINLESS STEEL CASTINGS

E.L. Brown, T.A. Whipple, and G. Krauss

Colorado School of Mines

Golden, Colorado 80401

ABSTRACT

The metallography of duplex stainless steel castings

was studied in a series of CF8M alloys with varying ferrite

content. Various etching techniques were employed to facilitate microstructural characterization of primary and

secondary structures by light microscopy. These observations were then correlated with transmission electron microscopy observations of substructure and selected area electron diffraction studies of the crystallography of the 6 y

transformation. All alloys possessed a significantly recovered dislocation substructure in both the y and 6 phases

which was correlated with certain microstrucutral features.

The crystallography of the 5 - y transformation was found to

be consistent with a Kurd jumov-Sachs type orientation relationship regardless of apparent residual 6 morphology. On

the basis of these observations a model for the evolution of

residual 6 morphology was proposed.

243

INTRODUCTION

The solidification behavior of austenitic stainless

steels is of concern in both welding and casting technology.

The term "austenitic" stainless steel is actually a misnomer

with respect to weldment and casting microstructures since

these can contain significant amounts of delta-ferrite (6),

the high temperature BCC phase of iron, as well as the FCC

austenite phase (y). The presence of 6 can significantly

alter the physical and mechanical properties of the y

matrix. Property variations due to 5 are beneficial or deleterious dependent upon the nature of the material application and its temperature regime. For example, in the case

of weldments the presence of 5-10% 6 has been reported to

reduce the incidence of hot cracking ( 1 , 2) and microfissuring.(3) At high temperatures (>700°C) 6 can transform to

the brittle sigma (a) phase, resulting in a degradation of

mechanical properties(A-6) , and in highly corrosive environments it can provide a path for corrosion attack. (7)

Ferrite in an austenite matrix also generally increases the

strength and lowers the fracture toughness of welds and

castings(8), this is especially the case in cryogenic applications. A general review of 6 in austenitic stainless

steel follows. This review will consider the effects of

245
composition, solidification and solid state transformation

on the extent and morphology of 6 in welds and castings.

Review of 6 in Austenitic Stainless Steel

Composition Effects . The overwhelming volume of work

has been conducted in the area of welding of austenitic

stainless steels and, more specifically, concerning the

affects of various filler metal compositions on 6 content.

Some additional work has been performed to extend the concepts employed in weldments to castings. The effects of

composition on 6 content are generally embodied in the form

of so-called Cr- and Ni -equivalents . The former integrate

the effects of ferrite formers and the latter the effects of

austenite formers. The Schaeffler and BeLong Diagrams

relate 6 content to Cr- and Ni-equivalents . (9-11) Similar

empirical correlations between 6 content and alloy composition have been established for castings and are contained in

the Schoefer Diagram. (12, 13) The data contained in these

diagrams refer to the 6 content at ambient temperature,

which formed in a non-equilibrium manner. In fact, the

equilibrium microstructures of these alloys at ambient

temperature generally do not contain any 6 . The departure

from equilibrium can be effected by altering the solidification rate, the cooling rate subsequent to solidification

and/or modifying the chemical composition. ( 14)

Solidification Effects . Numerous investigations have

been conducted to elucidate the solidification behavior of

this class of stainless steels. The solidification behavior

can be divided into primary 6 and primary y modes according

to the ratio of chromium and nickel equivalents. This

division is frequently illustrated by construction of appropriate constant Fe pseudo-binary sections of the Fe-Cr-Ni

ternary . (15,16,25) It is important to realize that pseudobinary sections through ternaries are only qualitative. (2-

6) They are, however, useful for qualitatively Illustrating

solidification modes as a function of composition. A typical pseudo-binary of this sort is depicted in Figure 1.

The two-phase, y + L and 6 + L regions are separated by

a eutectic triangle of y + 6 + L region. Below the triangle

a y + 6 phase field exists. It can be seen that If the

nominal composition of the steel possesses a low Cr

e^ (low

proportion of ferrite formers) the primary solid phase is

austenite. Similarly, in a steel with a high Cr (high

246
Primary y

Solidification

Primary 8

Solidification

Figure 1. Constant Fe pseudobinary phase diagram

proportion of ferrite formers) ferrite is the primary solid

phase. The solidification mode becomes more complex in the

vicinity of the eutectic triangle. From nominal compositions > Cr

e >

C the primary solid phase is still austenite, howevi?, ferrite precipitates before solidification

is complete. Likewise, for nominal compositions > Ni >

C

2

the primary solid phase is ferrite but austenite precipitates prior to complete solif ication. As will be discussed

later, the microstructures that result from solidification

and cooling in the vicinity of the eutectic triangle are also more complex. This finer classification of solidification modes has been described in the literature. (14,17-

20,25,27) The ternary eutectic behavior described above is

applicable to the majority of commercial austenitic stainless steel grades that contain < 90 w/o iron. In the materials under consideration, the eutectic is of a divorced

nature since the amount of liquid at the eutectic composition is generally small. In primary y solidified alloys the

divorced phase at the subgrain and grain boundaries is 6

.

In primary 6 solidified alloys the divorced phase

is y.(21) Studies of solidification in conventional ingots,

complemented by a directional solidification and quenching

technique, have confirmed the potentially dynamic nature of

the solidification sequence of primary 6 alloys. (14,16,

247
25,27,28) Initial precipitation of 6 could be interrupted

by precipitation of y, in the form of an envelope, before

complete solidification.

Concomitant with the evolution of a solidification

structure is the redistribution of solute elements, commonly

referred to as solute segregation. The various possibilities have been described in detail elsewhere. (29-31) Of

course, equilibrium solidification is not a reasonable

assumption. A more realistic situation is one in which

solid state diffusion is minimal so that coring or microsegregation occurs during solidification. (30,31) Perfect

mixing in the liquid and mixing in the liquid by diffusion

only represent extreme situations. Real cases probably

exist between these two extremes. For the cases described

above^the final distributions of solute after normal freezing

are depicted in Figure 2.

A generalized theory of primary 6 solidification of

austenitic stainless steel has been formulated for weldments, assuming mixing in the liquid by diffusion only. (15)

The transients are associated with the subgrain (dendrite)

core and the subgrain boundary (interdendritic) region.

Basically, Cr is segregated to the dendrite core and Ni to

the interdendritic region. Attempts have been made to

experimentally determine solute distribution via electron

microprobe and scanning transmission electron microscopy

studies. (15, 16, 18-20, 22, 23, 27, 28) Difficulties arise in

separating the solute redistribution due to solidification

and that due to the solid state 6 y transformation, without making assumptions regarding the nature of this phase

transformation, e.g., no solute redistribution during 6 -» y

transformation. (15, 2 2, 23)

Solid State 6 ->• y Transformation and Residual

Morphologies in Primary 6 Alloys. The 6 content of as- solidified primary 6 alloys can be substantially reduced

during cooling and post solidification treatments. (24,32)

Therefore, it is obvious that a 6 -* y transformation plays

a major role in the evolution of microstructure in duplex

austenitic stainless steel welds and castings.

A 6 + y transformation only occurs for alloys that

undergo primary 6 solidification and the nature of the

transformation appears to vary with composition, cooling

rate and perhaps with solidification behavior. It is important to note that a number of labels have been associated

with the various observed residual morphologies that are

248
descriptive in nature but may not accurately describe the source of the morphology.

Primary 6 alloys remote from the eutectic triangle

solidify with almost 100 v/o 6 ; y forming interdendritcally. In this type of alloy the resultant room temperature

microstructure is completely determined by the solid state

transformation. The character of the phase transition

Figure 2 Schematic diagram of solute distribution

after normal freezing.

depends essentially on the degree of undercooling, which in

turn, depends upon the cooling rate. (14,15) For relatively

rapid cooling the 6 y transformation has been alternatively referred to as a Widmanstatten or an acicular( 17,20)

morphology. The Widmanstatten phase transformation can be

described as a thermally activated, and diffusion controlled

reaction in which an orientation relationship between the

parent and product phases is established. (33) The residual 6 product has been termed "lathy”. (17, 18, 20) Microstructural investigation of the lathy morphology in

welds(18) has indicated the presence of several differently

oriented y-6 lath bundles in a given large growth unit.

Furthermore, the residual 6 is not related to the primary

6 dendritic structure, due to solidification. (14,18,25) At

DISTANCE

249
slower cooling rates the mode of <5 + y transformation in

these types of alloys has been termed "equiaxial" ( 17 , 20) and

it is not exactly clear what type of 6 morphology this

results in.

The majority of primary 6 alloys of interest intersect

the eutectic triangle so that y envelopes the primary

6 prior to completion of solidification. The solid y envelope allows for heterogeneous nucleation of y during the

6 + y transformation at smaller undercooling (higher temperature). Therefore, the resultant 6 exhibits less of the

Widmanstatten type morphology. (14, 16, 17, 20, 25, 28) The 6 morphology that results in this case has been termed "vermicular" because of its so-called roundish appearance. (15,20,

21) A continuous spectrum of ferrite morphology, from

vermicular to lathy, has been observed with increasing

Cr /Ni .(20) Regardless of the observed morphology,

the4

6 in primary 6 alloys is definitely located at the cores

or spines of dendrite arms (subgrains) .(14,16,17,20,21,25)

The progressive growth of y into the primary 6 results in

segregation of "f erritizers” to the 6 at the dendrite or subgrain core, thus stabilizing it against transformation.

(10,21,22,28) The dendritic form is now easily

defined. (14, 25) Implicit in this observation is that formation of a y envelope occurs with attendant segregation of

alloying elements. Since diffusion, and hence homogenization of alloying elements is more difficult in austenite,

the segregation pattern persists and allows observation of

the prior dendritic form at room temperature . (25) In alloys

which solidify as vertually 100 v/o 6 the ease of diffusion

in this phase minimizes segregation at high temperature and

therefore the dendritic growth form is not observable.

Limited transmission electron microscopy (TEM) observations have shown that lathy 6 possesses a Kurd jumov-Sachs

(K-S) type orientation relationship(34) with austenite,

whereas vermicular 6 possesses no crystallographic

relationship with austenite. (20) Angular faceting of

6/y interfaces In primary 6 alloys has also been observed

and noted in castings . (21) The observations of

primary 6 alloys described above have all implied a diffusion controlled mechanism for the 6 -* Y

transformation. ( 16,20)

Some of the observations described above have been

synthesized to rationalize the behavior of larger castings

and ingots. Castings frequently exhibit a columnar-to250
equiaxed grain transition during solidification, i.e.,

columnar grains growing inwards from the mold wall and a central equiaxed grain zone. In alloys that solidify almost

completely as 6 the first y precipitates in the primary

grain boundaries and intragraunularly as the temperature

drops. All of these regions of y have different orientations and the columnar-to-equiaxed transition is thus

observable in the primary (solidification) and secondary

(solid state) microstructures . (25) In alloys that Intersect

the eutectic triangle the columnar-to-equiaxed grain transition is masked by the 6 + y transformation. Austenite

begins to solidify from the melt after initial 6 solidification and prior to completion of solidification.

Columnar y crystals nucleate at high temperature and grow

under steep temperature gradients across the entire ingot

cross-section. With the high driving force for y growth

thus produced, changes in the original 6 orientation present

no barrier to that growth. (19, 28, 35) Mixtures of these two

modes of behavior are frequently observed In the same ingot. Once 5 has solidified, the y envelope necessary

for the L-HS+L+S+y transition is difficult to form at high

growth rates. (25) Therefore, completely 6 solidification

might be expected in the vicinity of the mold wall and a mixed mode in the ingot interior. Variations in 6 level

through an ingot cross-section must be rationalized in light

of the fact that higher 6 levels, associated with higher

cooling rates, can be offset by rapid homogenization in the

finer solidification structures .(24,32) Rapid

homogenization lowers the 6 content. As the 6 levels in

welds and castings increase the <$ phase becomes more

continuous . (37) Furthermore, upon dissolution of 6,

spheroidization can take place concurrently . (24)

The objective of the present investigation was to

characterize the microstructure and substructure of CF8M

duplex stainless steel castings. Effort was made to correlate microstructural aspects such as 6 morphology, observed

via light microscopy, with substructure, observed by transmission electron microscopy. In addition, the crystallography of the 6 - y transformation was studied by selected

area electron diffraction. Structural and crystallographic

data were synthesized and a model for the evolution of 6 morphology proposed.

251
EXPERIMENTAL PROCEDURES

Chemistries and Processing of Alloys

A number of CF8M alloy castings were produced to ASTM

specification A351 . The alloy designations, their chemical

compositions and ferrite content are included in Table I,

below. The ferrite contents shown were determined by quantitative metallography techniques. Alloys K-9, K-14 and K24 were all produced separately as keel block castings

(approx. 20 cm x 8 cm x 14 cm) whereas the other alloys were

produced together as a series of controlled chemistry castings In which the ferrite content was purposely varied at a constant interstitial content (C+N). This material was produced In the form of fairly large castings (10 cm x 30 cm

x 60 cm). The alloy production and castings procedures have

been described by Finch. (38) Subsequent to casting, all

material was annealed at approximately 1400 K.

TABLE I Chemical Composition

Alloy

Deslg-

£ N Mn SI Cr Ni Mo v/o

6-ferrite

nation

K-9 0.02 _ 1.06 1.10 19.33 11.67 2.14 9 K-14 0.05 - 0.83 1.25 19.50 12.20 2.69 14 K-24 0.07 - 0.52 1.32 21.48 9.48 3.18 24

1 0.06 0.05 1.10 0.59 18.04 13.24 2.12 1.1

2 0.06 0.05 0.30 1.16 19.62 10.61 2.12 8.2

3 0.06 0.05 0.36 1.13 19.23 8.25 2.16 17.2

4 0.06 0.05 0.35 0,95 20.40 8.22 2.13 23.1

5 0.05 0.05 0.39 1.11 22.54 9.84 2.26 28.5

Light Microscopy

Specimens for light microscopy were ground and lapped

to a metallographic finish by standard procedures. Four

etching techniques were employed. A deposition color etch

composed of 100 ml distilled b^O, 20 ml HC1 and 2 g potassium metabisulfite was applied as a light etch to delineate

the prior solidification structure and as a heavy etch to

delineate the macrostructure of the castings. As a light

etch, a blue interference layer is deposited on the regions

of the original 6 dendrites (high Cr) and a gold layer on

the interdendritic regions (low Cr). The latter actually

represent regions of Interdendritic liquid. An electro-etch

252
of 10 g oxalic acid in 100 ml distilled i^O (2-6 V and room

temperature) was employed to delineate boundaries between

5 and y and twin boundaries. When more contrast was required between 6 and y a ION KOH electrodeposition etch

was employed to deposit a light blue to light brown film on

the 6 ferrite.

Transmission Electron Microscopy (TEM)

Thin foils were produced for TEM by cutting wafers ~ 0.4 mm in thickness with a low speed diamond saw. The wafers

were reduced to ~ 0.07 mm in thickness by mechanically

abrading and 3 mm diameter disks produced by slurry

drilling. The latter is a low temperature-low deformation

procedure. Foils were produced by electropolishing to

perforation in a Fischione twin jet polisher, employing a 5v/o HC10^ -25v/o glycerin-ethanol electrolyte (40 V- 80 mA

at room temperature). Foils were examined in a Philips EM

400 transmission electron microscope, operated at 120 kV.

RESULTS

Light Microscopy

Some aspects of the solidification macrostructure are illustrated in the K-9 alloy, light color etched, Fig.

3(a). A columnar-to-equiaxed grain transition occurred

characterized by an increase in interdendritic liquid and

detatched columnar and equiaxed dentrites. The residual

6 was seen to exist as networks at the dendrite spines,

Fig. 3(b). When a heavier color etch was applied to the same casting section, a coarse columnar grain structure

parallel to the direction of maximum heat flow was observed,

Fig. 3(c). This columnar grain structure appeared to grow

across the previously described columnar-to-equiaxed grain

transition. A higher magnification macrograph shows that

the dendritic structure, as defined by the residual 6,

possesses the same orientations across many of the columnar

grain boundaries. Fig. 4. In addition, the coarse columnar

structure appears composed of a finer scale mosaic of different contrast regions that are faceted on a microscale.

The specimens from alloys 1-5 were obtained from regions

corresponding the central equiaxed region of the castings.

The residual 6 morphologies in these alloys are generally

illustrated in Fig. 5. As the ferrite content is increased

253
Figure 3(a) Columnar-to-equiaxed grain morphology

transition in solidification structure.

Figure 3(b) Residual 6 at spines of dendrite arms.

254
Figure 3(c) Columnar y crystals formed

during 5+y phase transformation

the ferrite microstructure becomes coarser and exhibits

greater continuity. On a given plane of polish the 6 in the

higher ferrite alloys appears composed of either relatively

long, filamentary particles or equiaxed cell-like

networks. There are also many faceted 6/y interfaces. A

facet is defined as a straight seqment of 5/y interphase

boundary. Frequently, facets seem to occur at definite,

recurring angles to one another on a given plane of polish,

suggestive of facet variants. Twinning of the y matrix was

also observed, Fig. 6. The 6/y interface morphology is

quite intricate and can generally be described as facets

connecting curved seqments of interface. Closer examination

of the interfaces revealed the presence of steps on faceted

and curved segments, Fig. 7.

Transmission Electron Microscopy

General Features of Substructure . A well recovered

dislocation substructure was observed in all castings.

Although no quantitative measurements were made, it appeared

that the subgrain structure was coarser in the 1.1% 6 specimen in comparison to those containing higher amounts of

ferrite, Fig. 8.

255
Figure 4 Columnar y crystal structure

showing residual 5 at spines of dendrite

arms and variations in contrast.

Figure 5(a) Residual 6 morphology as a function

of 6 content, (a) 1.1% 6

256
Figure 5(b) 9% 6

Figure 5(c) 28.5% 6

257
Figure 6 Twinning in y matrix

Figure 7 Residual 6 showing ledges on

S/y interface.

258
Figure 8(a) Recovered dislocation substructure

TEM (a) 1.1% 6

Figure 8(b) 28.5% 6

259
The subgrain boundaries were sometimes composed of dislocation tangles, seemingly "knit" together, while in other

instances a well defined dislocation net characterized the

boundary, Fig. 9.

Figure 9(b) Weak Beam CDF

260
The ferrite also contained a substantial amount of dislocation substructure. In fact, the 28.5% 6 exhibited a "bamboo" subgrain structure in the ferrite. Fig. 10.

Sub-boundaries were observed to bridge the width of the ferrite region.

Figure 10 Recovered dislocation substructure

in 5, 28.5% 6; TEM.

In addition to subgrains there were also significant

levels of individual dislocations in the subgrain interiors,

apparently in the form of debris. The dislocations in the

1.1% 6 specimen were frequently extended, as evidenced by

the extended nodes in the sub-boundary of Fig. 9. There was

less evidence of extended dislocations in specimens containing ferrite. There was also more evidence of planar

slip in the 1.1% 6 specimen despite the presence of a recovered subgrain structure, Fig. 11.

Large populations of stacking and twin faults were

observed. The diffraction contrast from faulted regions was frequently indicative of overlapping stacking faults, Fig

12(a). These faults were observed to emanate from subgrain

boundaries and 5/y interfaces. There was also evidence of

lamellar twins that were associated with discrete

reflections in diffraction patterns Fig. 12(b). Their

presence is consistent with the observations of light

microscopy. Faults were sometimes associated with edge

261
Figure 11 Evidence of planar slip in y, 1.1% 6; TEM

Figure 12(a) Faulting in y; TEM

(a) Overlapping stacking faults

262
Figure 12(b) Twin faults.

cracks in the foil. It is difficult to state whether these

cracks are actually present in the bulk material or are an artifact of foil making.

Crystallography of the 6 » y Transformation . The

crystallography of the 6 - y transformation was studied

extensively by TEM of thin foils obtained from the keel

block casting containing 9 v/o 5. However, the features

described below were also observed in specimens containing

higher ferrite contents. Observations of the

6/y interfacial morphology indicated the presence of many

facets and steps. No attempt was made to exactly determine

the lattice correspondence between austenite and ferrite.

However, SAD evidence was consistent with the widespread

occurrence of Kurdjumov-Sachs (K-S) - type orientation relationships, Fig. 13. Curved and straight interfaces were examined and the lattice correspondences described above

persisted regardless of the extent of boundary curvature.

Single surface trace analysis of straight sections of

interface indicated habit planes slightly deviated

from {ill} Y . Small facets on the 6/y

interface were observed and there was some evidence of

misfit dislocation structures in the boundaries, Fig. 14.

*The Kurdjumov-Sachs, Nishayama-Wasserman and Bain orientation relationships, relating bcc and fee phases, are very

similar.

263
Figure 13(a) 6/y Interphase interface, TEM (a) BF

[TT i]

s -- [To.]

y

Figure 13(b) SAD consistent with K-S

orientation relationship

264
Figure 14(a) 6/y Interphase interface, TEM

(a) Interfacial dislocations

Figure 14(b) Boundary facets.

265
DISCUSSION

These CF8M alloys, with the possible exception of the

1.1% 6 alloy, solidify with 6 as the primary solid phase,

Fig. 3(b). The color etching illustrates that a significant

portion of the 6 that forms upon solidification subsequently transforms to y during cooling, via a solid state

6 + y transformation. The two macrostructures of Figs.

3(a) and (c) provide important information in further

characterizing the transformation sequence. The completely

columnar secondary structure of Fig. 3(c) does not exhibit

a columnar-to-equiaxed grain transition and therefore must

be associated with the solid state 6 -* y transformation. (35) The fact that dendrites are observed to grow

across the columnar grain boundaries of the secondary structure also supports the assertion that this structure is

produced by the 6 - y transformation, Fig. 4. Furthermore,

observation of the secondary structure is consistent with

the following transformation sequence: L+L+S+L+S+y^S+Y

(14,25) In this case a y envelope begins to solidify from

the melt and columnar y crystals can nucleate

heterogeneously at high temperature and grow with the aid of

very steep temperature gradients. Once nucleated under

these conditions, a columnar y crystal will grow right across mlsorientations between 6-ferrite dendrites and

parallel to the path of maximum heat flow. It is possible

that since the secondary Y structure was not observed in the

higher 6-ferrite content castings (>20v/o), these alloys

possess the following transformation sequence:

L-»-L+6>6-»-6+y •

The significant levels and the recovered nature of

dislocation substructure might be expected to derive from

several sources. Strain can occur due to solidification

shrinkage, general thermal strains, differential thermal

expansion and transformation strains. It Is probable that

the total strain induced in the material from all these sources is sufficient to account for the dislocation densities observed. The thermal activation attendant with

annealing and cooling is sufficient to recover the dislocation substructure through subgrain formation. The observation of mosaic regions differing in contrast in Fig. 4 might

then be rationalized as manifesting an underlying subgrain

structure. The evolution of dislocation substructure could

also derive from a compromise between the tendency of the y

to grow along a direction of maximum heat flow and its

266
tendency to grow more readily along certain crystallographic

directions. The two directions are not necessarily parallel. One possible way to effect the compromise and accomodate the mismatch is the generation of dislocations. The

succeeding discussion on lattice correspondences lends

support to the notion of crystallographic directions of easy

growth.

The 6/y interfacial morphology, viewed via light and

transmission microscopy, and diffraction evidence supports

the existence of a lattice correspondence between 6 and y

regardless of the curvature of the interface between the two

phases. This lattice correspondence is embodied in a K-Stype orientation relationship ({lll}y//{110}6 :'<110>tf//<lll><5)

that is associated with a high degree of lattice matching

across certain interfaces. If the 6/y boundary is oriented

favorably a high degree of lattice matching will result in a low specific interfacial energy and, therefore, a facet. In

the case of a K-S type lattice correspondence a high degree

of atomic matching exists parallel to {111} y and {110} 6 in

the direction of <110> y and <111> 6 but for only several

atomic distances from an arbitrary atomic origin. However,

good matching resumes on the next lower or upper

{111} y: {110} 6 pair. The resulting interfacial structure may be thought of as a series of steps or terraces

.

(39)

The macro habit will then be irrational, i.e., not exactly

parallel to {111} y , as was observed in the present investigation. In addition, the interfacial dislocation structures frequently observed in this investigation are generated to accomodate losses in coherency due to lattice

mismatch and changes in boundary orientation.

The existence of a K-S-type orientation relationship

and a faceted 6/y interfacial morphology suggests a general

model for the 6 - y transformation anologous to those proposed for diffusion controlled formation of Widmanstattentype proeutectoid a. (40-43) A faceted nucleus of y forms

either at solidification subgrain boundaries or on the y

envelope formed by direct y solidification from the melt.

These alternative modes of y nucleation would not alter the

basic morphological evolution. A low energy, faceted interface on the nucleus minimizes total interfacial energy and

therefore the activation barrier for nucleation. Growth of

nucleii is then accomplished through the movement of curved,

incoherent segments of 6/y interface. These possess high

energy and high mobility compared to the low mobility of the

faceted portions of interface. Movement of faceted bound267
aries can occur via movement of interface ledges or steps

parallel to the interface. Steps on 6/y interfaces have

been viewed by light and transmission electron microscopy

(Figs. 7 and 14). The features of the residual 6 microstructure can then be viewed as evolving through the rapid

movement of incoherent segments of 6/y interface with

coherent segments being dragged along behind. The coherent

segments move more slowly via a ledge mechanism. A schematic diagram showing the evolution of the residual 6 morphology is shown in Fig. 15. This schematic emphasizes the

continuous nature of the 6 network and that the appearance

of filamentary particles or cell-like networks on a particular plane of polish represent sectioning effects.

Figure 15 Schematic diagram of evolution of residual

6 morphology via 6+y transformation

CONCLUSIONS

(1) All of the alloy castings possessed a recovered dislocation substructure in both the 6 and y phases. A

recovered dislocation substructure was rationalized on

the basis of induced strains due to solidification, thermal expansion, transformation and differential thermal

expansion.

268
( 2 ) Selected area electron diffraction evidence confirmed

the widespread occurrence of a K-S type orientation

relationship between 6 and y . This prevailed across

curved and faceted sections of 6/y interface. Data on

the faceted segments were consistent with a habit plane

which deviated slightly from {llljy.

(3) In an alloy in which y precipitation occurred before

completion of solidification it was shown that the

differences between y growth along preferred crystallographic directions and in the direction of maximum heat

flux might be accomodated by formation of a recovered

dislocation substructure.

(4) The evolution of residual 6 morphology may be considered as occurring via the rapid movement of incoherent (curved) segments and the relatively slow movement

of semicoherent (faceted) segments by a diffusion

controlled process. This is accomplished while maintaining a lattice correspondence in a manner similar to

the formation of Widmanstatten proeutectoid ferrite

from austenite.

REFERENCES

1. H.F . Reid and W.T. Delong, Met. Prog., Vol. 84 (1973),

p. 73.

2. H. Fredriks and L.J. VanderTorn, Br. Weld. J. , Vol. 15

(1968). p. 178.

3. C.D. Lundin and D.F. Spond, Weld. J. Res. Suppl. Vol.

55 (1976), p. 356-s

.

4. R.T. King, N.C. Cole, and G.M. Goodwin, J. Pressure

Vessel Technol., Vol. 98 (1976), p. 194.

5. L.R. Poole; Met. Prog., Vol. 65 (1954), p. 108.

6. O.H. Henry, M.A. Cordovi, and G.J. Fisher, Weld. J.

Res. Suppl., Vol. 34 (1955), p. 75-s.

7. K.E. Pinnow and A. Moskowits, Weld. J. Res. Suppl.,

Vol. 49 (1970), p. 278-s.

8. V.N. Zemsin and G.L. Petrov, Svar. Proisvod., No. 5

(1967). pp. 6-8.

9. W.T. DeLong, Weld. J. Res. Suppl., Vo. 53 (1974), p.

273-s

.

10. A. Schaeffler, Metal Progress * Vo. 56, (1949), pp. 680-

680B.

11. C. Long and W.T. DeLong; Weld. J., Res. Suppl, Vol. 52

(1973), pp. 281-s-297-s

.

12. F.H. Beck, E.A. Schoefer, J.W. Flowers, and M.G.

Fontana, in Advances in the Technology of Stainless

Steels and Related Alloys, ASTM STP 369, (1965), pp.

159-174.

269
13. L.J. Schwartzendruber, L.H. Bennett, E.A. Schoefer,

W.T. DeLong, and H.C. Campbell, Weld. J. Res. Suppl.,

Vol. 53 d(1974) , p. 1-s.

14. J. Beech: Personal Communication.

15. J.C. Lippold and W.F. Savage; Weld. J. Res. Suppl.,

Vol. 58 (1979), p. 330-s.

16. S.A. David, G.M. Goodwing, and D.N. Braski, Weld. J.

Res. Suppl., Vol. 58 (1979), p. 330.

17. N. Suutala, T. Takalo, and T. Moisio, Met. Trans. A,

Vol. 10A (1979), p. 512.

18. N. Suutala, T. Takola, and T. Moisio, Met. Trans. A,

Vol. 10A (1979), p. 1183.

19. T. Takalo, N. Suutala, and T. Moisio, Met. Trans. A,

Vol. 10A (1979), p. 1173.

20. N. Suutala, T. Takato, and T. Moisio, Met. Trans. A,

Vol. 11A (1980) p. 717.

'

21. M.J. Cieslak and W.F. Savage, Weld. J. Res. Suppl.,

Vol. 59 (1980), p. 136-s.

22. J.C. Lippold and W.F. Savage; Weld. J. Res. Suppl.,

Vol. 59 (1980), p. 48-s

23. C.E. Lyman, Weld. J. Res. Suppl., Vol. 58 (1979), p.

189-S.

24. V.S. Raghunathan, V. Seetharaman, S. Venkadesan, and P.

Rodriquez, Mt. Trans A, Vol. 10A (1979), p. 1683.

25. J.M. McTighe and J. Beech, J. Res., SCRATA, No. 37,

(1977), p. 43.

26. G. Masing, Ternary Systems: Introduction to The Theory

of Three Component Systems, Dover, New York, (1960).

27. 0. Hammar and U. Svensson, Solidification and Casting

of Metals, Proceedings of an International Conference

on Solidification, University of Sheffield, England

(1977), p. 401.

28* H. Fredriksson, Met. Trans., Vol. 3 (1972), p. 2989.

29. J.W. Rutter and B. Chalmers; Can. J. Phys., Vol. 31

(1973) p. 15.

30. J.W. Christian, The Theory of Tranformatlons in Metals

and Alloys , Pergamon Press, New York (1965), p. 566.

31. W.A. Tiller, K.A. Jackson, J.W. Rutter, and B.

Shalmers, Acta Metall . , Vol. 1 (1953), p. 428.

32. Y. Kinoshita, S. Tadeda, and H. Yoshimura, Tesu-toHagane, Vol. 65 (1979), pp. 1176-1185.

33. J.B. Newkirk, Precipitation from Solid Solution ASM

Seminar, American Society for Metals, (Nov. 1957).

34. G. Kurdjumov and G. Sachs, Ann. Phys., Vol. 64 (1930),

p. 325.

35. A. Hultgren, J.I.S.I.Vol. 70 (1929), pp. 69-113.

270
36. T.P.S. Gill, R.K. Dayal, and J.B. Gnanomoorthy , Weld.

J. Res. Suppl., Vol. 58 (1979), p. 375-s.

37. R.G. Berggren, N.C. Cole, G.M. Goodwin, J.O. Stiegler,

G.M. Slaughter, R.J. Gray, and R.T. King, Weld. J. Res.

Suppl. , Vol. 57 (1979), p. 167-s.

38. L. Finch, Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures, NBSIR-1645, (1981),

p. 337.

39. M.G. Hall, H.I. Aaronson, and K.R. Kingman, Surface

Science, Vol. 31, (1972), pp. 257-274.

40. W.C. Johnson, C.L. White, P.E. Marth, P.K. Ruf, S.M.

Tuominen, K.D. Wade, K.C. Russell, and H.I. Aaronson,

Met. Trans. A , Vol. 6A (1975), p. 911.

41. H.I. Aaronson, Decomposition of Austenite by Diffusions! Processes , The Metal Society, American Institute

of Mining, Metallurgical, and Petroleum Engineers, New

York (1962), p. 387.

42. H.I. Aaronson, C. Laird, and K.R. Kinsman, Phase Transformations , ASM seminar, American Society for Metals,

Metals Park, Ohio, (October 1968) p. 313.

43. H.I. Aaronson, J.K. Lee, and K.C. Russell, Precipitation Processes in Solids , The Metal Society, American

Institute of Mining, Metallurgical and Petroleum

Engineers, New York (1976).

271

A TRANSMISSION ELECTRON MICROSCOPY STUDY OF CAST DUPLEX AUSTENITIC STAINLESS STEELS

WITH VARYING FERRITE CONTENT: AS-CAST AND DEFORMED AT 4 K

E. L. Brown

Department of Metallurgical Engineering

Colorado School of Mines

Golden, Colorado 880401

ABSTRACT

The evolution of dislocation substructure and residual 6 morphology in the as-cast state was studied by transmission electron microscopy and selected area electron diffraction in a series of

CF8M alloys with varying ferrite content. Results were correlated with a previous light microscopy

study. A well recovered dislocation substructure existed in both phases and was correlated with

the 6+y transformation and thermal history. The lattice correspondence observed by electron diffraction and habit planes of facets in the 6/y interfacial morphology were consistent with a model

based on a semicoherent interface. All results were consistent with a proposed model for evolution

of residual 6 morphology analogous to the proeutectoid y+a transformation. The evolution of substructure was detailed in 17.2 volume percent and 28.5 volume percent alloys that were strained at

4 K. Low strain specimens were characterized by dislocation generation in both phases in the vicinity of the 6/y interface, dislocation-dislocation interactions and dislocation-sub-boundary

interactions in y. Later stages displayed formation and intersection of shear/deformation bands in

Y and was followed by a" formation at the intersections. Bands of dislocations formed in the 6 followed by deformation twinning in that phase. The results of y deformation were considered in

light of current theories of the strain- induced y+a' transformation and work hardening in metastable y in addition to theories regarding plastically non-homogeneous alloys.

273

INTRODUCTION

The solidification behavior of 300-series austenitic stainless steels is of concern in both

welding and casting technology. In the cast or welded state these alloys can contain up to approximately 30 volume percent delta-ferrite (6), the high temperature BCC phase of iron. For this reason they are often referred to as duplex stainless steels. The presence of 6 can significantly

alter the physical and mechanical properties of the austenite (y) matrix. Property variations due

to 6 can be beneficial or deleterious dependent upon the nature of the material application and its

temperature regime. For example, in the case of weldments the presence of 5-10 volume percent 6 has been reported to reduce the incidence of hot cracking [1,2] and microfissuring |_3]. At high

temperatures (>700°C) 6 can transform to the brittle sigma phase, resulting in a degradation of

mechanical properties [4-6], and in highly corrosive environments it can provide a path for cor- rosion attack [7]. Ferrite in an austenite matrix also generally increases the strength and lowers

the fracture toughness of welds and castings L8].

Due to the high strength and toughness of wrought austenitic stainless steels at cryogenic

temperatures [9-11], these alloys are currently being considered for applications in the con- struction of superconducting magnets as both weld filler metals and castings. There is some con- cern over the usefulness of ferrite containing alloys in cryogenic applications since at these

temperatures the ferrite is below its ductile-to-brittle transition. Studies of the strength and

fracture toughness of 300-series welds at 4 K indicated the same general trend, cited above, of

increasing strength and decreasing toughness with increasing 6 content [12,13], There has been

considerably less work done to characterize the deformation and fracture behavior of duplex castings at cryogenic temperatures [14-16].

The amount and morphology of 6 in a weld or casting has been found to depend upon composition,

solidification behavior, and a solid state 6+y phase transformation. The effects of these factors

on the extent and morphology of residual 6 has been extensively investigated and reviewed [17-20].

These alloys can generally be subdivided according to their solidification behavior into primary 6 and primary y alloys dependent upon which is the leading phase on solidification. This is determined almost completely by alloy composition. Primary 6 alloys contain higher levels of ferrite

stabilizers, such as chromium, and primary y alloys contain higher levels of austenite stabilizers,

such as nickel. The leading phase and solute redistribution have been found to determine the mor- phology and location of ferrite in the post-solidification microstructure. In primary y alloys

enrichment of the intercellular or interdendritic regions in chromium stabilizes ferrite in these

regions. In primary 6 alloys chromium is segregated to the dendrite or cell cores during solidification and 6->- y transformation. Therefore, ferrite in primary 6 alloys is located at dendrite

spines. The morphology of the 6 at dendrite spines in primary 6 alloys varies from vermicular or roundish, to lathy, to acicular appearing as the nominal alloy composition varies from positions

within the eutectic triangle (low contents of ferrite stabilizers) to positions substantially re- moved from the eutectic triangle (high contents of ferrite stabilizers). Mechanisms to rationalize

the evolution of residual ferrite morphology have been formulated on the basis of light and trans- mission electron microscopy (TEM) [17,19]. However, the majority of TEM studies to date have been

mainly concerned with high resolution x-ray microanalysis to determine the nature of solute gradients after solidification.

275
In an effort to systematically study the effects of ferrite and interstitial content on deformation and fracture of castings at 4 K, a series of controlled chemistry castings were produced and

tested. The light microscopy of these alloys has been described previously [17,20] and is characterized by well defined facets on 6/y interfaces and regularly occurring angles between facet vari- ants. The deformation and fracture behavior of these castings at 4 K have also been described

elsewhere [15,16]. Uniaxial tensile tests were performed in which strain was monitored from the

microstrain region to fracture and were characterized by three stages of work hardening, similar to wrought 300-series steels tested at that temperature. Deformation twinning in 6 and martensitic

transformations in y were observed in the microstructures of strained specimens. The work hardening behavior was found to be dependent upon ferrite content. Fracture toughness was also found to vary with ferrite content in a manner similar to weldments.

Based on the similarities in microstructure and mechanical behavior between welds and castings

it may be possible to extend conclusions drawn from microstructural observations of castings to welds. This would greatly facilitate modelling of weld behavior since structural observations in welds are more difficult to accomplish due to their finer microstructure.

The objective of this paper is to describe TEM studies performed to characterize the substructure of as-cast and deformed specimens of cast duplex stainless steel. Effort was made to correlate microstructural aspects such as 6 morphology, observed previously by light microscopy,

with the crystal lography of the 6+y transformation , studied by selected area electron diffraction.

The validity of a previously proposed model for the evolution of residual 6 morphology was investigated. The formation of substructure in 6 and y, with strain at 4 K, was observed in addition to the nature of strain- induced martensite (a') formation in y. Deformation modes in 6 were also

described.

EXPERIMENTAL PROCEDURES

A number of CF8M alloy castings were produced to ASTM specification A351. The alloy designations, their chemical compositions and ferrite content are included in Table I. The ferrite

contents shown were determined by quantitative metal 1 ographic techniques. Alloys K-9, K-14, and

K-24 were all produced separately as keel block castings (approx. 20 cm x 8 cm x 14 cm) whereas the

other alloys were produced together as a series of controlled chemistry castings in which the ferrite content was purposely \raried at a constant interstitial content (C+N). The controlled chemistry castings were relatively large (10 cm x 30 cm x 60 cm). The alloy casting procedures have been

described by Finch [21]. Subsequent to casting, all material was annealed at approximately 1400 K.

Mechanical Testing

The details of tensile testing at 4 K have been described previously [15,16]. Specimens of

Alloys 3 and 5, strained to levels of 0.005, 0.02, and fracture were sectioned for TEM observation.

276
Microscopy

Specimen preparation for light microscopy has been described by Brown [17]. Thin foils were produced for TEM by cutting wafers ^ 0.4 mm in thickness with a low speed diamond saw. Wafers from

the tensile specimens were removed from the gage section perpendicular to the tensile axis. The

wafers from specimens strained to fracture were removed from regions of uniform strain. Wafers were reduced to ^ 0.07 mm in thickness by mechanically abrading and 3 mm diameter disks produced by

slurry drilling. The latter is a low temperature, low deformation procedure. Foils were produced

by electropolishing to perforation in a Fischione twin jet polisher, employing a 5 volume percent

perchloric acid -25 volume percent glycerin-ethanol electrolyte (40 V, 80 mA at room temperature)

.

Foils were examined in a Philips EM 400 transmission electron microscope, operated at 120 kV. Scanning electron microscopy (SEM) and energy dispersive X-ray analysis (EDS) were performed on selected specimens for the purpose of determining the extent of alloying element partitioning between 6 and y.

RESULTS

Undeformed Specimens

The microstructures of duplex austenitic stainless castings that solidify in a primary 6 mode

have been studied by light microscopy [17,20]. Relevant observations from that work are extracted

here. The residual 6 morphologies were characterized by faceted 6/y interfaces. Faceted segments

of 6/y interface frequently occurred at definite angles with respect to one another on a given

plane of polish and high magnification light microscopy has revealed ledges on straight and curved

segments. Fig. 1. The habit planes of facet variants could not be determined unambiguously from

measurements of inter- variant angles on a given plane of polish since the crystallographic orientation of the plane of polish was not generally known. However, the approximate grouping of intervariant angles suggested that the habits might be in the vicinity of {110)6. Furthermore, the

observation of facets and ledges suggested some sort of lattice correspondence between 6 and y, namely a Kurdjumov-Sachs (K-S) type orientation relationship.

Another pertinent metallographic result from previous studies regards the observation of

columnar y crystals due to the 6->y transformation. The origins of columnar y crystals have been

discussed in detail [17,20]. These crystals grow along directions of maximum heat flow and closer

examination reveals subtle differences in contrast between regions inside the crystals and dentate

boundaries between these regions, Fig. 2. The boundary structure within the columnar y crystals is reminiscent of subgrain structures observed by other investigators with the aid of polarized light

microscopy [22].

A well recovered dislocation substructure in y was observed in all castings, Fig. 3. The

subgrain boundaries were sometimes composed of dislocation tangles, Fig. 4. In other instances a well defined dislocation net characterized the boundary, Fig. 5. Note the extended nodes in the

boundary net of Fig. 5, indicative of a low stacking fault energy material. Measurements of the

subgrain size, by a linear intercept method, indicated an average subgrain size of 2-5 y and no consistent variation with ferrite content. The subgrains also varied in shape from equiaxed to

277
somewhat elongated. The long segments of elongated subgrains were generally attached to 6/y interfaces and stretched across the y separating two regions of 6. This observation was more prevalent

in the higher 6 content alloys. The ferrite also contained a substantial amount of dislocation

substructure. In fact, Alloy 5 (28.5 volume percent 6) exhibited a "bamboo" subgrain structure in the ferrite, Fig. 6. The subgrain boundaries in y and 6 can produce varying degrees of cusping in the 6/y interface where they intersect it, Fig. 7. Another type of y dislocation network was observed in these alloys that was generally characterized by straight segments of network attached to the 6/y interface, Fig. 8. The planes of these networks were consistent with { 1 1

1

} y and { 1 10 } y . Similar networks have been observed in lightly deformed austenitic stainless steel and rationalized

in terms of interactions between glide and forest dislocations at pile-ups in the vicinity of grain

boundaries [23,24]. The presence of { 1 10 }y networks may be explained by the fact that { 1 10}y is one of the types of planes of "easy climb" in fee metals and alloys [25]. Networks are formed by a process called "knitting."

In addition to subgrains in both phases, there were also significant levels of dislocations,

in the form of debris. The debris is characterized by individual dislocations, pile-ups, and

tangles, indicative of dislocation interactions. Dislocation debris is formed in 6 and y. The

dislocations are sometimes associated with subboundaries, Fig. 9(a). The dislocations may have

been produced at the subboundary or just interacted with it after production elsewhere. Dislocation debris was also observed in association with 6/y interfaces, Fig. 6, 9(b). These dislocations

were produced at the interface. The tangles of dislocations in 6 were sometimes very dense and

associated with heavily jogged dislocation lines and loops, 9(c). The dislocations in 6, indicated

at the arrow in Fig. 6, seem to be at an earlier stage of interaction. Dislocations emitted from

the subboundary at left are interacting with forrest dislocations. Jogs and dipoles are the re- sult. Overlapping stacking faults were occasionally observed to emanate from 6/y interfaces, Fig.

6. It is uncertain whether these are characteristic of the substructure or artifact.

The crystallography of the 6->y transformation was studied extensively in Alloy K-9 by TEM and

selected area electron diffraction (SAD). However, the features described below were also observed

in specimens containing higher ferrite contents. Observations of the 6/y interfacial morphology

indicated the presence of many relatively large facets and steps. No attempt was made to exactly

determine the lattice correspondence between 6 and y. However, numerous SAD observations were

found to be consistent with the widespread occurrence of Kurdjumov-Sachs (K-S)-type orientation

relationships, Fig. 10.* Curved and straight segments of interphase interface were examined and

the lattice correspondences described above persisted regardless of the extent of boundary curvature. Single surface trace analysis yielded a population of habit plane trace normals that ap- peared to cluster in the vicinity of ( 1 1 2 ) y , on the boundary of the stereo triangle and near

(011)y, Fig. 11.

*The Kurdjumov-Sachs, Ni shayama-Wasserman and Bain orientation relationships, relating bcc and fee

phases, are related to one another by minor rotations.

278
This data illustrates the irrational nature of the habit plane for "straight" segments of

interphase interface. Small steps were observed on the 6/y interface, Fig. 12(a). The 6/y interface is further characterized by an interfacial substructure, Fig. 12(b).

Deformed Specimens

Specimens of Alloys 3 and 5 (17.2 and 28.5 volume percent 6) strained to 0.005, .02 and fracture at 4 K were observed by TEM. The evolution of dislocation substructure and second phase (

a

') formation with strain was similar for the two alloys; therefore representative observations are presented for the strain sequence.

At a strain of 0.005 dislocation sources in y are activated in the 6/y interface. Dislocations on intersecting slip planes were observed to interact to produce tangled networks. Fig. 13. Note the complicated contrast of the boundary associated with activation of dislocation sources,

Figs. 13(c), (d). Dislocations can actually be observed to bow out from the sources in the 6/y

interphase interface (arrow in Fig. 13(c), (d)). Dislocation sources were also observed in the

vicinity of y subboundary intersections with the 6/y interface, Fig. 14. The 6/y interface is cusped at this location. These are likely to be regions of high stress concentration due to the

overlapping stress fields of dislocations and the torques imposed on the 6/y interface. The dislocation density in y seems to be higher in the vicinity of the 6/y interface, where tangles of dislocations were frequently observed. Fig. 15. Glide dislocations have also interacted with subboundaries from the as-cast state, Fig. 16. Actually, 6 dislocation sources are also activated, as can

be seen in Fig. 17, where dislocations are observed to bow out into both phases from the 6/y interface.

The structures of specimens strained to 0.02 were characterized by intersections of deformation or shear bands, Fig. 18. The deformation bands appear to be composed of a higher density of

dislocations from intersecting slip planes then observed at lower strain. Some of the other bands

appear to be associated with stacking fault fringe contrast. These might contain e-martensite,

although none was detected by electron diffraction. Martensite has begun to form at the intersections of deformation and shear bands. Fig. 19. There is considerable accommodation strain

associated with a" formation as evidenced by dislocations generated in the vicinity of intersections with a " , Fig. 20. The a' thus formed sometimes possessed projected images polyhedral in morphology. Fig. 19(a). Most of the a.' formed at intersections possessed a single habit variant.

However, there were some regions in which bands of a' formed, containing multiple a' habit var- iants, Fig. 21. The bands are bundles of faults bounded by { 1 1 1 }y . In the 6, dislocation pile-ups

were observed at the 6/y interfaces. Fig. 22. These dislocations probably emanated from the interface. The dislocations in the pile-ups are jogged and there is a mottled contrast that appears

crystallographically arranged (at arrow in Fig. 22). Since the trace of this contrast is consistent with {110)6 it is probably due to dislocations viewed end-on. The jogs may be due to interactions with the end-on dislocations. This strain level also appeared to signal the onset of

deformation twinning in the 6, Fig. 23. Note the kink at the 6/y interface induced by twin

(arrow)

.

279
Specimens strained to fracture were characterized by a greater number of shear/deformation

band intersections and, therefore, a' formation. The 17.2 volume percent 6 alloy (Alloy 3) displayed regions, sometimes in close proximity, of widely varying degrees of a' formation, Fig. 24. The upper region of y contains a" forming at shear/deformation band intersections. Fig. 24(b). The

lower y region has been almost' totally transformed to a". Fig 24(c). Selected area electron diffraction of the a.' in this region indicated variants that are twin related. There may also be fine

twins in plates of a'. The 28.5 volume percent 6 alloy (Alloy 5) was not generally characterized

by as much a' formation. However, this alloy did possess an increased number of deformation/shear

band intersections with a' formation when compared with specimens of the same alloy at lower

strains, Figs. 25, 26. Deformation substructure in the 6 also seems to have evolved further with

increased strain. The bands formed by intersecting dislocations have now extended across the re- gions of 6, Fig. 27. In addition, macroscopic deformation twins have been observed by light microscopy and TEM, Fig. 27a, Fig. 28. Deformation has occurred ahead of growing twins to accommodate

their growth, Fig. 29(a), and in the y where a twin abuts the 6/y interface, Fig. 29(b). It ap- pears that dislocations induced by accommodation strains have interacted with dislocations from the

bands in 6 to form tangles. Fig. 29(a).

It should be noted that all SAD evidence on a' was consistent with a K-S orientation relationship between y and a'. Also single surface trace analyses of a' habit planes were consistent with

a { 225 }y habit, within the obvious errors inherent in this kind of analysis.

Energy Dispersive X-ray Analysis

Energy dispersive x-ray analyses (EDS) performed on these alloys indicated partitioning of

alloying elements. Chromium and molybdenum segregate to 6 and nickel to y, Fig. 30. A number of

EDS spectra were obtained from 6 and y in Alloys 3 and 5 and subjected to a standardless ZAF semiquantitative analysis. The results, depicted in Fig. 31, are not to be construed as fully quantitative but rather to display trends in alloying elements that affect austenite stability.*

Chromium and nickel both increase in y as the volume percent 6 increases from 17.2 to 28.5.

DISCUSSION

Undeformed Specimens

These CF8M alloys, with -the possible expection of Alloy 1 (1.1 volume percent 6), solidify

with 6 as the primary phase. The micrographs of Fig. 1 represent residual 6 that remains after a 6+y transformation . The columnar y crystals observed in Fig. 2 are also the result of this 6+y

transformation and this is supported by the observation that the dendrite structure is continuous

across the columnar grain boundaries, Fig. 2(b). The dendrite structure is defined by the residual

6 at dendrite spines.

Comparison of low magnification scans, that would indicate overall chemical composition, with the

results of analytical chemical analysis (Table 1) indicate small discrepancies

.

280
The observations of Rhines [22] indicate that the variations in contrast and dentate boundaries observed in Fig. 2 could be due to a mosaic or subgrain structure. The plausibility of this

assertion was confirmed by TEM observations of well recovered subgrain structures. Fig. 3-6. The

strains necessary to produce these subgrain structures might be derived from several sources. Strain can occur due to solidification shrinkage, strains developed during cooling of a large

casting in which parts of the casting cool at varying rates, differential thermal expansion or contraction of different phases [& and y) and transformation (6-*y). The last two sources have been

considered in some detail [26]. It was determined from dilatometry that a significant amount of

strain was imparted to the material by the 6+y transformation . It is probable that the total

strain induced from all these sources is sufficient to account for the observed dislocation densities. The thermal activation attendant with annealing and cooling is sufficient to recover the

dislocation substructure through subgrain formation. Other investigators have observed the formation of subgrains during solidification and for solid state transformation [25] and cite the fact

that conservative climb of dislocations aided by pipe diffusion of vacancies can produce subgrain

boundary dislocation networks very rapidly. This process is termed "knitting" and can produce all

of the boundary networks observed. The evolution of a subgrain structure could also be viewed as deriving from a compromise between the tendency of the y to grow along a direction of maximum heat

flow and its tendency to grow more readily along certain crystallographic directions, i.e. to maintain a lattice correspondence between <5 and y. The two growth directions are not necessarily par- allel. The mismatch could be accommodated by the generation of subgrain structure. The succeeding

discussion on lattice correspondence lends support to the notion of crystal lographic directions of

easy growth.

The 6/y interphase interfacial morphology, viewed via light and transmission electron microscopy, supports the existence of a lattice correspondence between 6 and y. Electron diffraction

evidence, in fact, confirms the existence of a K-S-type orientation relationship between 6 and y, Fig. 10. The variability in the lattice correspondence has been cited by Brooks, et. al. [19] who

found the lattice correspondence between 6 and y in welds to vary from the Kurdjumov-Sachs to the

Nishayama-Wasserman orientation relationship. The common feature of these types of orientation

relationship is the parallelism between closest packed planes in the two phases, i.e. { 1 10 } 6 //

{ 1 1

1

} y . This would seem to 'suggest a { 1 1 1 >y habit for the straight segments of 6/y’interface but

these segments of interface were observed to deviate substantially from { 1 1 1 }y . The deviation in

habit for fcc:bcc boundaries has been investigated in the past [27-32]. All of these investigations emphasize the partially coherent nature of the interface for all lattice correspondences.

Close matching across { 1 1 1 } y // { 1 10 } 6 is actually confined to relatively small regions. The area

of close matching in the interface is increased substantially by the creation of monoatomic ledges,

the broad faces of which contain the regions of good fit on the next lower pair of { 1 10 } 6 //

{ 1 1 1 } y . These are called structural ledges and account for the deviation in habit from { 1 1 1 } y

.

The regions on the broad faces of structural ledges in between the regions of good fit can be made

partially coherent by the introduction of misfit dislocations. Therefore the 6/y interface can be

viewed as the superposition of arrays of structural ledges and misfit dislocations. It should be

noted that these structural ledges and misfit dislocations are sessile and therefore do not participate in growth mechanisms. The efficacy of this model in describing the 6/y interfaces being

considered in this study is demonstrated by the fact that the predictions of theory [27], for a similar lattice parameter ratio of the two phases, are similar to the habit planes observed for

281
faceted 6/y interfaces, Fig. 11. The observation of interfacial substructure on 6/y interfaces

also supports the notion of partial coherency. Fig. 12. The large steps on 6/y interfaces, observed via light microscopy (Fig. 1(b)) and TEM (Fig. 12(a)) must then be the growth ledges pro- posed by other investigators [27,29,30].

The existence of a K-S-type orientation relationship and a faceted 6/y interfacial morphology

suggests that the general model for the 6+y transformation previously proposed [17] is essentially

an accurate description. That model proposed a transformation mechanism analogous to that of the

diffusion controlled formation of proeutectoid a from y. The 6->-y transformation has been established as diffusion controlled [18,19] so that the model would appear to apply in that regard. The

production of "lathy" 6 in austenitic welds due to a 6-»y transformation has been rationalized on the basis of an interface instability [19]. Closer scrutiny may reveal a mechanism equivalent to that described here.

Deformed Specimens

The stress-strain (o-e) behavior at 4 K of Alloys 1-5 have been determined [15,16], The o-e curves of all the alloys displayed three stages of work hardening behavior, typical of metastable

steels that undergo a y-*a'transformation [33-36]. The deformed specimens observed in this study were strained into Stage I (0.005), the beginning of Stage II (0.02) and Stage III (strain to fracture). It should be noted that the elongation was greater for the 17.2 volume percent 6 alloy than

for the 28.5 volume percent 6 alloy (19% vs. 11%) but the ultimate tensile strengths were almost

equal (1143 MPa vs 1164 MPa). The TEM work in this investigation confirmed that no a' forms in Stage I, as noted by other workers [35,36]. This stage was associated with the generation of dislocations, deformation bands and shear bands. It was interesting that no direct evidence of

e-martensite formation could be obtained via electron diffraction. The volume fraction of e should

increase to a maximum and then gradually decrease with strain [33,36]. The maximum of (30-40

volume percent) occurs in AISI 304-L stainless steel at a strain of approximately 0.03 at T = 20°K

[33]. The maximum is expected to be less in magnitude and occur at lower strains at 4 K. Therefore e would not be expected in specimens strained to fracture. The maximum in e occurs at the

onset of the transition from Stage II to Stage III (low to high rate of work hardening) [36]. It

is possible that compositional effects could have suppressed formation of e by their effect on stacking fault energy. This possibility was not evaluated. An experimental reason for difficulty

in detecting e via electron diffraction was the unfavorable y textures introduced by preferred

crystallographic growth during solidification [19]. Specimens of all alloys were characterized by

a high propensity for <100>y crystal lographic zones near the foil normal. This is not a favorable

austenite zone to detect e reflections since there is no low order e zone axis parallel to a <100>y

zone axis. This is due to the prevailing lattice correspondence between y and e-martensite

[37,38].

The dislocations produced in the vicinity of the 6/y interface during this stage of

deformation (Fig. 15) may be associated with production of a population of geometrically neces- sary dislocations [15,39]. Geometrical ly necessary dislocations are generated by the microstructure in response to non-homogeneous strain i.e. matrix strain gradients in the presence of

hard particles. A statistically stored dislocation density is produced by random dislocation

282
interactions. Whipple [15] assumed that at the outset of deformation the geometrical ly necessary

dislocation density would be more significant than the statistically stored density and that the

former should increase with volume fraction 6. Therefore, the initial work hardening rate should

increase with volume fraction 6. This was observed. However, in the present study dislocations

were observed to tangle with one another and interact with as-grown subboundaries. Fig. 13 and 16. These would seem to be processes for statistical storage of dislocations. In addition, it is un- certain what effect a 2-4 ym diameter subgrain structure has on the initial work hardening rate. Previously produced subgrain structures have been observed to increase the flow stress of a ma- terial in a manner functionally similar to the well-known Hall-Petch grain size strengthening

effect [40-43]. In fact, Sherby [41] found that wel 1 - recovered subgrain boundaries were more effective barriers to dislocation motion than subboundaries formed at lower temperatures. The

additional dislocation sources at the 6/y interface that were activated in the presence of abutting

subgrain boundaries might also be expected to increase the effect of subboundaries on initial work

hardening, Fig. 14. It would seem that although a model based on initial generation of geometrically necessary dislocations might still be an accurate description, the transition to control of

deformation behavior by statistically stored dislocations is somewhat nebulous, at least at the

lowest level of strain observed in this study. It is probably not controlled by the difference

between "austenite grain size" and spacing, as proposed by Whipple [15]. The variation in friction

stress, T , identified by Whipple [15], may have an additional origin in the as-grown subgrain

structure.

The work hardening characteristics of these alloys are also affected by a' formation in Stage II. This is not accounted for by Ashby's model of plastically inhomogeneous alloys. The a' formed in Stage II nucleates at shear and deformation band intersections. Figs. 19 and 20. The

observations described concur with current theories of strain-induced a' formation [35,37,38,44,

47-52]. Intersections are favorable sites for heterogeneous nucleation of a.' because of the high

stress concentrations present there and because the necessary dislocation configurations for nu- cleation of the bcc phase are also present. According to the available literature a" nucleated

during straining possesses a lath morphology that extends along the length of the intersection, a K-S or Ni shayama-Wasserman lattice correspondence, a single habit of either { 225 > y or { 259 }y and is bounded by { 1 1 1 } y of one of the intersecting bands. The laths grow via coalescence after repeated

nucleation with straining [52]. The coalesced morphologies resemble polyhedra and will grow out of

the bounding { 1 1 1 }y . Similar morphologies have been observed in this study. Fig. 19 and 20. There

were regions of multiple habit variant a'. Fig. 21. According to Reed [36], multiple habit variant

a' is associated with a.' that formed spontaneously on cooling. This is surprising in view of the

fact that a

is supposed to be virtually suppressed on quenching to 4 K [45,47].

Formation of a' signals the onset of a low work hardening region. Some investigators have

proposed that the a' formed at this stage is non-hardening and provides "windows" for slip [35].

They cite a special variant of the K-S orientation relationship for the first formed laths. This

variant results in virtual matching of a' and y dislocation Burgers vectors.

The fractions of a.' increased dramatically for both alloys in Stage III and so did the work

hardening rate. It is interesting that a' promotes low work hardening rates in Stage II and high

work hardening rates in Stage III [3,6]. It is proposed that higher work hardening results from a'

283
formation in active slip bands and the fact that e martensite levels decrease is cited as supporting evidence. The number of deformation/shear band intersections also increased during this

stage of deformation, lending support to the model for a' nucleation kinetics proposed by Olson and

Cohen [49]. The twin related a' variants observed Fig. 24(c) have been observed by other workers

[46]. These alloys only appeared to differ in the levels of a" after straining to fracture. The

28.5 volume percent 6 alloy contained less a". This is puzzling since the work hardening rates were apparently equal. However, the flow stress in the 28.5 volume percent 6 alloy was higher and

the elongation was less. It is possible that necking and failure occurred sooner for this alloy

and therefore, less a' was formed. A problem arises with this explanation in that, qualitatively,

the number of deformation/shear band intersections appeared to be the same for both alloys. The

analysis of Olson and Cohen [49] therefore suggests that a diminished thermodynamic driving force

may be responsible for less a' formation. The EDS measurements suggest increasing nickel and

chromium levels in the y as volume fraction 6 increases, Fig. 31. A rough calculation of the

driving force for the y+a' transformation, aFy "*“

, was made employing an equation due to Breedis

and Kaufman [53] for Fe-Cr-Ni alloys. The calculation indicated higher austenite stability for the

28.5 volume percent 6 alloy. The austenite stability in duplex alloys must therefore be considered

from the standpoint of alloy partitioning.

It is uncertain what effect the as-grown subgrain structure had on a' formation. Cellular

dislocation substructures that raise the strength of austenite, were observed to suppress a' formation [54]. The potential effects of a subgrain structure on work hardening and a' formation will

have to be evaluated more fully.

In the modelling of deformation behavior in these alloys the ferrite has been assumed to be

non-deforming. However, the TEM observations in the present investigation would suggest appreciable deformation during straining, although some of the dislocation substructure could have been

produced by residual strains introduced during specimen reheating to room temperature after

testing. The deformation twinning observed in ferrite actually requires prior deformation for

nucleation of twins [55]. Twins apparently nucleate at dislocation pile-ups. It is possible that

the dislocation pile-ups observed in the 6 provide nucleating sites for twins, Figs. 22 and 27. Emissary slip has been proposed as a mechanism for accommodating the high stresses in the vicinity

of the tip of a twin [55]. This is observed in Fig. 29(a). The high stress at the tip of a twin

that intersects the 6/y interface is accommodated by intense deformation in the adjacent y, Fig. 29(b). It is uncertain whether deformation twinning affects mechanical behavior directly or what implications the prior deformation in 6 has for modelling. However, twinning has been asso- ciated with the origin of cleavage nucleii in a-iron [56]. In this model interactions between

emissary dislocations are proposed to induce the formation of crack nucleii. Therefore, it -would

appear that deformation twinning in 6 could promote fracture via cleavage cracking in 6 or failure

of the y in the vicinity of a 6/y interface where a twin has intersected it to produce intense

accommodation strain.

284
CONCLUSIONS

(1) Residual 6 morphology is consistent with a diffusion controlled mechanism analogous to that

proposed for the proeutectoid y+ct transformation. The habit plane of faceted segments of 6/y

interface and a Kurdjumov-Sachs-type lattice correspondence are both consistent with predictions based on a semicoherent interphase interface model.

(2) A well recovered subgrain structure was observed in 6 and y for all alloys in the as-cast

state. A well recovered dislocation substructure was rationalized on the basis of induced

strains due to solidification, thermal contraction, differential thermal contraction and 6+y

transformation

.

(3) TEM observations indicated that low strain behavior was characterized by dislocation production in the vicinity of 6/y interfaces, dislocation-dislocation interactions and dislocationsubboundary interactions. Later stages of deformation were characterized by intersecting

shear/deformation band intersections and a' formation at the intersections. These observations were correlated with three stage hardening mechanisms. Differences in a' formation in various alloys could be due to variations in austenite stability attendant with alloy partitioning.

(4) Ferrite was found to deform prior to deformation twinning.

RECOMMENDATIONS FOR FUTURE WORK

(1) Transmission microscopy investigations of the dislocation and ledge nature of semicoherent 6/y

interfaces. This would be a valuable fundamental contribution to the understanding of this

generally significant interface type. The 6/y interfacial structure is probably a very significant factor affecting the ultimate 6 morphology in welds and castings.

(2) More extensive studies of the evolution of dislocation substructure and a' formation during

straining over a wider’ range of strain and ferrite content. These could be correlated with

existing theories regarding work hardening and the y+a' transformation . TEM observations

should be supplemented with some sort of bulk detection techniques for e and a', such as x-ray

or magnetic measurements. The latter could be made in situ.

(3) Investigate more fully the nature of deformation in 6, deformation twinning in 6 during

straining at 4K and the implications of twinning on the deformation and fracture behavior of

welds and castings.

285
REFERENCES

[1] H. F. Reid and W. T. Delong: "Making Sense Out of Ferrite Requirements in Welding Stainless

Steel." Met. Prog ., Vol. 84 (1973), p. 73.

[2] H. Fredriks and L. J. VanderToorn: "Flot Cracking in Austenitic Stainless Steel Weld

Deposits," Br. Weld. J ., Vol. 15 (1968), p. 178.

[3] C. D. Lundin and D. F. Spond: "The Nature and Morphology of Fissures in Austenitic Stainless

Steel Weld Metals" Weld J. Res. Suppl . Vol. 55 (1976), p. 356-s.

[4] R. T. King, N. C. Cole, and G. M. Goodwin, "Creep Behavior of SMA Type 316 Weld Metal with

Controlled Residual Elements," J. Pressure Vessel Techno! ., Vol. 98 (1976), p. 194.

[5] L. R. Poole: "An Unwanted Constituent in Stainless Weld Metal," Met. Prog., Vol. 65 (1954),

p. 108.

[6] 0. H. Henry, M. A. Cordovi , and G. J. Fischer, "Sigma Phase in Austenitic Stainless Steel

Weldmqnts," Weld. J. Res. Suppl ., Vol. 34 (1955), p. 75-s.

[7] K. E. Pinnow and A. Moskowitz: "The Corrosion Resistance of Stainless Steel Weldments," Weld.

J. Res. Suppl ., Vol. 49 (1970), p. 278-s.

[8] V. N. Zemzin and G. L Petrov: "Influence of Ferrite Content on the Properties of Austenitic

Weld Metal," Svar. Proizvod ., No. 5 (1967), pp. 6-8.

[9] R. P. Reed, F. R. Fickett, M. B. Kasen, and H. I. McHenry, "Magnetic Fusion Energy Low

Temperature Materials Program: A Survey," in Mat'ls Studies for Magnetic Fusion Energy

Applications at Low Temperatures - I , NBSIR 78-884, (1978), pp. 245-335.

[10] H. I. McHenry and R. P. Reed, "Structural Alloys for Superconducting Magnets in Fusion Energy

Systems," Nuclear Eng. Design , Vol. 58, (1980), pp. 219-236.

[11] R. L. Tobler, D. T. Read and R. P. Reed, "Strength and Toughness Relationship for

Interstitial ly Strengthened AISI 304 Stainless Steels at 4 K," in Mat'ls Studies for Magnetic

Fusion Energy Applications at Low Temperatures - IV , NBSIR 81-1645, (1981), pp. 37-75.

[12] D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr., "Metallurgical Factors

Affecting the Toughness of 316L SMA Weldments at Cryogenic Temperatures," in Mat'ls Studies

for Maqnetic Fusion Enerqy Applications at Low Temperatures - II, NBSIR 79-1609, (1979), pp. 313-352.

[13] E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA Austenitic Stainless Steel

Weld Metals: Part I - Role of Ferrite," Weld. J. Weld. Res Suppl. Vol 57, (1978), pp. 325s- 333s

.

[14] T. A. Whipple and H. I. McHenry, "Evaluation of Weldments and Castings for Liquid Helium

Service," in Mat'ls Studies for Magnetic Fusion Enerqy Applications at Low Temperatures - IV,

NBSIR 81-1645, (1981), pp. 273-287.

[15] T. A. Whipple, "The Effect of Delta-Ferrite on the Mechanical Properties of CF8M Stainless

Steel Castings at 4K," Mat'ls Studies for Maqnetic Fusion Energy Applications at Low

Temperatures - V , NBSIR 82-1667, (1982), pp. 289-317.

[16] T. A. Whipple and E. L. Brown, "Deformation and Fracture of Stainless Steel Castings and

Weldments at 4 K," in Trends in Welding Research in the U.S. , ASM Conference Proc., S. A. David, ed, New Orleans, La., Nov. 1981, pp. 601-621.

[17] E. L. Brown, "Metallography of 6-Ferrite in Austenitic Stainless Steel Castings," in Mat'ls

Studies for Maqnetic Fusion Energy Applications at Low Temperatures - IV, NBSIR 81-1645,

(1981), pp. 387-414.

[18] S. A. David, "Ferrite Morphology and Variations in Ferrite Content in Austenitic Stainless

Steel Welds," Weld J,, Weld. Res . Vol. 60, (1981), pp. 83s-71s.

286
[19] J. A. Brooks, J. C. Williams and A. W. Thompson, "Solidification of Austenitic Stainless

Steel Welds," in Trends in Welding Research in the U.S. , ASM Conference Proc., S. A. David,

ed, New Orleans, La., Nov 1981, pp. 331-357.

[20] E. L. Brown, T. A. Whipple and G. Krauss, "Metallography of Duplex Stainless Steel Castings,"

ASM sponsored symposium, St. Louis, Mo., Oct. 1982, (in press).

[21] L. Finch, "The Influence of Ferrite and Nitrogen on Mechanical Properties of CF8M Cast

Stainless steel," in Mat'ls for Magnetic Fusion Energy applications at Low Temperatures - IV,

NBSIR 81-1645, (1981) pp. 337-355.

[22] F. N. Rhines, "Evolution of Microstructure During Hot Plastic Deformation," Chapt. 10 in The

Inhomogeneity of Plastic Deformation , ASM, Seminar, Oct. 1971, pp. 251-284.

[23] M. J. Whelan, P. Hirsch, R. W. Horne, and W. Bollman, "Dislocations and Stacking Faults in

Stainless Steel," Proc. Roy. Soc. A , Vol . 240, (1957), pp. 524-538.

[24] M. J. Whelan, "Dislocation Interactions in Face-Centered Cubic Metals With Particular

Reference to Stainless Steel," Proc. Roy. Soc. A , Vol. 249, (1959), pp. 114-137.

[25] H. M. Miekk-oja and V. K. Lindroos, "The Formation of Dislocation Networks," Surface Sci.,

Vol. 31, (1972), pp. 422-455.

[26] J. W. Elmer, "Thermal Expansion Characteristics of Stainless Steel Welds," MS Thesis,

Colorado School of Mines, Dec. 1981.

[27] M. G. Hall, H. I. Aaronson, and K. R. Kinsman, "The Structure of Nearly Coherent fcc:bcc

Boundaries in a Cu-Cr Alloy," Surface Sci . , Vol. 31, (1972), pp. 257-274.

[28] K. C. Russell, M. G. Hall, K. R. Kinsman, and H. I. Aaronson, "The Nature of the Barrier to Growth at Partially Coherent FCC:BCC Boundaries," Met. Trans ., Vol. 5, (1974), pp. 1503-1505.

[29] J.M. Rigsbee and H. I. Aaronson, "A Computer Modelling Study of Partially Coherent FCC:BCC

Boundaries," Acta Met . Vol. 27, (1979), pp. 351-363.

[30] J. M. Rigsbee and H. I. Aaronson, "The Interfacial Structure of the Broad Faces of Ferrite

Plates," Acta Met ., Vol. 27, (1979), pp. 365-376.

[31] P. Kluge-Weiss and H. Gleiter, "Electron Microscopic Observations on the Structure of

Dislocations in Interphase Boundaries," Acta Met. , Vol. 26, (1978), pp. 117-121.

[32] G. B. Olson, "Fine Structure of Interphase Boundaries," Acta Met., Vol. 29, (1981),

pp. 1475-1484.

[33] C. J. Guntner and R. P. Reed, "The Effect of Experimental Variables, Including the

Martensitic Transformation, on the Low-Temperature Mechanical Properties of Austenitic

Stainless Steels," Trans. ASM , Vol. 55, (1962), pp. 399-449.

[34] Dieter Fahr, "Stress-and Strain- Induced Formation of Martensite and Its Effects on Strength

and Ductility of Metastable Austenitic Stainless Steels," Met. Trans., Vol. 2 (1971),

pp. 1883-1892.

[35] K. Suzuki, T. Hashimoto and M. Ichihara, "An Experimental Study of the Martensite Nucleation

and Growth in 18/8 Stainless Steel," Acta Met. , Vol. 25, (1977), pp. 1151-1162.

[36] R. P. Reed, "Martensitic Phase Transformations," Chapt. 9 in Materials at Low Temperatures , Reed and Clark, eds., ASM, June 1983.

[37] J. A. Venables, "The Martensite Transformation in Stainless Steel," Phil. Mag ., Vol. 7,

(1962), pp. 35-44.

[38] P. L. Mangonon and G. Thomas, "The Martensite Phases in 304 Stainless Steel," Met. Trans ., Vol. 1, (1970), pp. 1577-1586.

[39] M. F. Ashby, "The Deformation of Plastically Non-Homogeneous Alloys," in Strengthen!' ng Methods in Crystals , Halsted Press Div., John Wiley & Sons, New York, (1971), pp. 137-192.

287
[40]

D. Kuhlmann-Wi 1 sdorf , "A Critical Test on Theories of Work Hardening for the Case of Drawn

Iron Wire," Met. Trans ., Vol . 1, (1970), pp. 3173-3179.

[41] C. M. Young and 0. D. Sherby, "Subgrain Formation and Subgrain Boundary Strengthening in Iron-Based Materials," J. I .S. I . , (Sept. 1973), pp. 640-647.

[42] C. Schwink and E. Gottler, "Dislocation Interactions, Flow Stress and Initial Work Hardening

of Copper Single Crystals with [100] Axis Orientation, "Acta Met., Vol. 24, (1976), pp. 173-179.

[43] H. J. McQueen, "The Production and Utility of Recovered Dislocation Substructures," Met. Trans. A , Vol. 8A, (1977), pp. 807-824.

[44] B. Cina, "Effect of Cold Work on the y->a Transformation in Some Fe-Ni-Cr Alloys," J I S I

,

(1954), pp. 406-422.

[45] R. P. Reed, "The Spontaneous Martensitic Transformation in 18% Cr-8%Ni Steels," Acta Met.,

Vol . 10, (1962) , pp. 865-877.

[46] J. F. Breedis and W. D. Robertson, "The Martensitic Transformation in Single Crystals of

Fe-Cr-Ni Alloys," Acta Met . , Vol. 10, ( 1962), pp. 1077- 1088.

[47] R. Lagneborg, "The Martensite Transformation in 18%Cr-8%Ni Steels," Acta Met., Vol. 12,

(1964), pp. 823-843.

[48] F. Lecroisey and A. Pineau, "Martensitic Transformations Induced by Plastic Deformation in the Fe-Ni-Cr-C System," Met. Trans ., Vol. 3, (1972), pp. 387-396.

[49] G. B. Olson and M. Cohen, "Kinetics of Strain Induced Martensitic Nucleation," Met. Trans. A,

Vol. 6A, (1975), pp. 791-795.

[50] G. B. Olson and M. Cohen, "A General Mechanism of Martensitic Nucleation: Part II. FCC+BCC

and Other Martensitic Transformations," Met. Trans. A , Vol. 7A, (1976), pp. 1905-1914.

[51] S. S. Hecker, M. G. Stout, K. P. Staudhammer, and J. L. Smith, "Effects of Strain State and

Strain Rate on Deformation- Induced Transformation in 304 Stainless Steel: Part I. Magnetic

Measurements and Mechanical Behavior," Met. Trans. A , Vol. 13A, (1982), pp. 619-626.

[52] L. E. Murr, K. P. Staudhammer, and S. S. Hecker, "Effects of Strain State and Strain Rate on Deformation- Induced Transformation in 304 Stainless Steel: Part II. Microstructural Study,"

Met. Trans. A , Vol. 13A, (1982), pp. 627-635.

[53] J. F. Breedis and L. Kaufman, "Formation of HCP and BCC Phases in Austenitic Iron Alloys,"

Met. Trans ., Vol. 2, (1971), pp. 2359-2371.

[54] J. F. Breedis, "Influence of Dislocation Substructure on the Martensitic^ Transformation in Stainless Steel ," Acta Met . , Vol. 13, (1965), pp. 2359-2371.

[55] S. Mahajan and D. F. Williams, "Deformation Twinning in Metals and Alloys," Int'l Met. Rev ., Vol. 8, (1973), pp. 43-61.

[56] A. W. Sleeswyk, "Twinning and the Origin of Cleavage Nucleii in a-Iron," Acta Met., Vol. 10,

(1962), pp. 803-812.

288
TABLE I Chemical Composition

A1 1 oy

Designation C N

K-9 0.02 -

K- 14 0.05 -

K- 24 0.07 -

1 0.06 0.05

2 0.06 0.05

3 0.06 0.05

4 0.06 0.05

5 0.05 0.05

Mn Si Cr

1.06 1.10 19.33

0.83 1.25 19.50

0.52 1.32 21.48

1.10 0.59 18.04

0.30 1.16 19.62

0.36 1.13 19.23

0.35 0.95 20.40

0.39 1.11 22.54

Ni_ Mo

Vol ume Percent

6- ferrite

11.67 2.14 9 12.20 2.69 14 9.48 3.18 24

13.24 2.12 1.1

10.61 2.12 8.2

8.25 2.16 17.2

8.22 2.13 23.1

9.84 2.26 28.5

289
LIST OF FIGURES

[ 1 ].

[ 2 ].

[3].

[4].

[5].

[ 6 ].

[7]

.

[ 8

]

.

[91.

[ 10 ].

[ 111 .

[ 12 ].

Residual 6 morphology.

(a) . 9 volume percent 6, facet variants. Light micrograph, oxalic acid electroetch.

(b) . 24 volume percent 6, ledges on 6/y interface (arrows). Light micrograph, KOFI el ectroetch

.

(a) . 9 volume percent 6, columnar y crystals. Light micrograph, deposition color etch.

(b) . 9 volume percent 6, interiors of columnar y crystals. Light micrograph, deposition color etch.

Recovered dislocation substructures . Transmission electron micrographs.

(a) . 1.1 volume percent 6.

(b) . 17.2 volume percent 6.

(c) . 28.5 volume percent 6.

8.2 volume percent 6, low angle boundary composed of dislocation tangles.

Transmission electron micrographs.

(a) . Bright-field.

(b) . Weak-beam dark-field.

1.1 volume percent 6, dislocation nets in subgrain boundaries. Transmission electron micrographs.

(a) . Bright-field.

(b) . Weak-beam dark-field.

28.5 volume percent 6, recovered dislocation substructure in 6. Transmission electron micrograph .

28.5 volume percent 6, low angle boundary in 6 and cusp at boundary - 6/y interface (arrow).

Transmission electron micrograph.

9 volume percent 6, dislocation network attached to 6/y interface. Transmission electron

micrograph

.

Dislocation debris. Transmission electron micrographs.

(a) . 28.5 volume percent 6, dislocations associated with sub-boundary

.

(b) . 28.5 volume percent 6, dislocations in 6 at 6/y interface.

(c) . 28.5 volume percent 6, jogged dislocations and loops in 6.

(a) . 9 volume percent 6, 6/y interface. Transmission electron micrograph.

(b) . SAD consistent with K-S orientation relationship.

Stereo triangle illustrating the results of single surface trace analyses of the habits of

6/y facets.

9 volume percent 6, aspects of 6/y interfacial morphology and structure. Transmission electron micrographs.

(a) . Steps on interface (arrow).

(b) . Interfacial dislocation arrays.

290
[13]

. Activation of dislocation sources at 6/y interface and interaction of dislocations in vicinity of interface, e=0.005. Transmission electron micrographs.

(a) . 17.2 volume percent 6.

(b) . 28.5 volume percent 6.

(c) . 17.2 volume percent 6; bright-field and weak-beam dark-field.

(d) . 28.5 volume percent 6; network and dislocation sources at 6/y interface.

[14]

. 17.2 volume percent 6, e=0.005; dislocation source at 6/y interface - sub-boundary junction.

Transmission electron micrograph.

[15]

, 17.2 volume percent 6, e=0.005; higher dislocation density in y at 6/y interface. Transmission electron micrograph.

[16]

. 28.5 volume percent 6, e=0.005; interaction of dislocations with as-grown subgrain boundary.

Transmission electron micrograph.

[17]

. 17.2 volume percent 6, e=0.005; dislocation sources (arrows) in 6/y interface bowing dislocations into both phases. Transmission electron micrograph.

[18]

, e=0.02, shear/deformation band intersections in y. Transmission electron micrographs.

(a) . 17.2 volume percent 6.

(b) . 28.5 volume percent 6.

[19]

. e=0.02, a' formation (arrows) at shear/deformation band intersections in y. Transmission

electron micrographs.

(a) . 17.2 volume percent 6.

(b) . 28.5 volume percent 6.

[20]

. 17.2 volume percent 6, e=0.02; dislocations generated in vicinity of forming a'. Transmission electron micrograph.

[21]

. 17.2 volume percent 6, e=0.02; multi-variant a' in shear/deformation bands (arrows indicate

traces of two variants). Transmission electron micrographs.

(a) . Bright-field.

(b) . Dark-field employing a' reflection.

(c) . High magnification micrograph of a' in region of arrow in (a) showing substructure

in a' platelets and at y/a' interfaces. The band is composed of a bundle of

faul ts

.

[22]

. 28.5 volume percent 6, e=0.02; dislocation pile-ups 6 at 6/y interface. Transmission electron micrograph.

[23]

. 28.5 volume percent 6, e=0.02; twin formation in 6. Transmission electron micrograph.

[24]

. 17.2 volume percent 6, e f racture l transmission electron micrographs.

(a) . Low magnification micrograph.

(b) . High magnification micrograph of lower y region.

(c) . High magnification micrograph of upper y region.

291
[ 25 ].

[26].

[27]

.

[28]

.

[29]. [30]

.

[31]

.

28.5 volume percent 6, e

f rac j- ure

*> a

' formation at shear/deformation band intersections.

Transmission electron micrographs.

(a) . Bright-field.

28.5 (b) . Dark-field employing a' reflection (arrows indicate positions of 6/y interface).

volume percent 6, e

f racture » a

' formation at shear/deformation band intersections.

Transmission electron micrographs.

(a) . Bright-field.

(b) . Dark-field employing a' reflection.

(a) and (b). 28.5 volume percent 6, . ; dislocation bands in 6. Transmission electron micrographs.

28.5

volume percent 6, e f rac -(- ure

; twins in 6. Transmission electron micrographs.

(a) . Bright-field of twins T^ and T^.

(b) . Dark-field of T^ employing twin reflection.

(c) . Dark-field employing 6 matrix reflection.

28.5

volume percent 6, e f racture ’ dislocation substructure produced by accomodation strains

associated with 6 twinning. Transmission electron micrographs.

(a) . In 6 ahead of growing twin (at arrow in Fig. 27(a)).

(b) . In y ahead of twin (arrow).

EDS line scans across region of 6 for Cr, Ni, and Mo indicating partitioning of elements.

EDS determination of Ni/Cr vs 6 content.

292
(a). 9 volume percent 6, facet variants. Light micrograph,

oxalic acid electroetch.

(b). 24 volume percent 6, ledges on 6/y interface (arrows).

Light micrograph, KOH electroetch.

Fig. 1. Residual 6 morphology

293
Fig. 2(a). 9 volume percent 6, columnar y crystals.

Light micrograph, deposition color etch.

294
Fig. 2(b). 9 volume percent 6, interiors of columnar y crystals.

Light micrograph, deposition color etch.

295
Fig. 3(a). 1.1 volume percent 6.

Fig. 3(b). 17.2 volume percent 6.

296
(c). 28.5 volume percent 6.

Fig. 3. Recovered dislocation substructures.

Transmission electron micrographs.

297
(b). Weak-beam dark-field.

Fig. 4. 8.2 volume percent 6, low angle boundary composed of

dislocation tangles. Transmission electron micrographs.

298
Fig. 5(a). Bright-field.

299
(b). Weak-beam dark-field.

Fig. 5. 1.1 volume percent 6, dislocation nets

Transmission electron micrographs.

in subgrain boundaries.

C

300
Fig. 6. 28.5 volume percent 6, recovered dislocation substructure in 6.

Transmission electron micrograph.

301
Fig. 7. 28.5 volume percent 6, low angle boundary in 6 and cusp at

boundary - 6/y interface (arrow). Transmission electron

micrograph.

302
Fig. 8. 9 volume percent 6, dislocation network attached to

6/y interface. Transmission electron micrograph.

303
Fig. 9(a). 28.5 volume percent 6, dislocations associated with sub-boundary

.

Fig. 9(b). 28.5 volume percent 6, dislocations in 6 at 6/y interface.

304
(c). 28.5 volume percent 6, jogged dislocations and loops in 6

Fig. 9. Dislocation debris. Transmission electron micrographs

305
Fig. 10(a). 9 volume percent 6, 6/y interface.

Transmission electron micrograph.

(IIO)

8 # (0 ZO)y

[11 -- [To

Fig. 10(b). SAD consistent with K-S orientation

].

],

relationship.

306
Fig. 11. Stereo triangle illustrating the results of single surface

trace analyses of the habits of 6/y facets.

307
Fig. 12(a). Steps on interface (arrow).

308
(b). Interfacial dislocation arrays.

Fig. 12. 9 volume percent 6, aspects of 6/y interfacial morphology

and structure. Transmission electron micrographs.

309
Fig. 13(a). 17.2 volume percent 6.

Fig. 13(b). 28.5 volume percent 6.

310
Fig. 13(c). 17.2 volume percent 6; bright-field and weak-beam dark-field.

311
(d). 28.5 volume percent 6; network and dislocation sources at 6/y interface.

Fig. 13. Activation of dislocation sources at 6/y interface

and interaction of dislocations in vicinity of interface,

e=0.005. Transmission electron micrographs.

312
Fig. 14. 17.2 volume percent 6, e=0.005; dislocation source at 6/y interface - sub-boundary junction. Transmission electron micrograph.

313
Fig. 15. 17.2 volume percent 6, e=0.005; higher dislocation density in

Y at 6/y interface. Transmission electron micrograph.

314
Fig. 16. 28.5 volume percent 6, e=0.005; interaction of dislocations with

as-grown subgrain boundary. Transmission electron micrograph.

315
Fig. 17. 17.2 volume percent 6, £=0.005; dislocation sources (arrows) in 6/y

interface bowing dislocations into both phases. Transmission electron

micrograph

.

316
Fig. 18(a). 17.2 volume percent 6.

317
(b). 28.5 volume percent 6.

Fig. 18. e=Q.02, shear/deformation band intersections in y. Transmission

electron micrographs.

318
Fig. 19(a). 17.2 volume percent 6.

319
(b) 28.5 volume percent 6.

Fig. 19. e=0.02, a' formation (arrows) at shear/deformation band intersections in

y . Transmission electron micrographs.

320
Fig. 20. 17.2 volume percent 6, e=0.02; dislocations generated in vicinity of

forming a'. Transmission electron micrograph.

321

V
Fig. 21(a). Bright-field.

1.0ft

Fig. 21(b). Dark-field employing a' reflection.

322
(c). High magnification micrograph of a' in region of arrow in (a) showing

substructure in a' platelets and at y/a' interfaces. The band is

composed of a bundle of faults.

Fig. 21. 17.2 volume percent 6, e=0.02; mul ti-variant a' in shear/deformation

bands (arrows indicate traces of two variants). Transmission electron

micrographs.

V

323
Fig. 22. 28.5 volume percent 6, e=0.02; dislocation pile-ups in 6 at 6/y

interface. Transmission electron micrograph.

324
Fig. 23. 28.5 volume percent 6, e=0.02; twin formation in 6. Transmission

electron micrograph.

V

325
Fig. 24(a). Low magnification micrograph.

326
Fig. 24(b). High magnification micrograph of lower y region.

Fig. 24(c). High magnification micrograph of upper y region.

Fig. 24. 17.2 volume percent 6, E f rac t ure

» transmission electron micrographs,

327
Fig. 25(a). Bright-field.

328
(b). Dark-field employing a' reflection (arrows indicate

positions of 6/y interface).

Fig. 25. 28.5 volume percent 6, E f rac t ure

; a

' formation at shear/deformation

band intersections. Transmission electron micrographs.

V

329
Fig. 26(a). Bright-field.

330
(b). Dark-field employing a' reflection.

Fig. 26. 28.5 volume percent 6, e f racture i a

' formation at shear/deformation

band intersections. Transmission electron micrographs.

331
(b)

Fig. 27. 28.5 volume percent 6, e f

. ; dislocation bands in 6. Transmission

electron micrographs. ra u

332
Fig. 28(a). Bright-field of twins Tj and T

2

.

Fig. 28(b). Dark-field of Tj employing twin reflection.

333
(c). Dark-field employing 6 matrix reflection.

Fig. 28. 28.5 volume percent 6, e f

. ; twins in 6. Transmission

electron micrographs. r

334
Fig. 29(a). In 6 ahead of growing twin (at arrow in Fig. 27(a)).

335
(b). In y ahead of twin (arrow).

Fig. 29. 28.5 volume percent 6, e f raC |- ure

i dislocation substructure produced by

accomodation strains associated with 6 twinning. Transmission electron

micrographs.

336
Fig. 30. EDS line scans across region of 6 for Cr, Ni, and Mo indicating

partitioning of elements.

337
Ni/Cr

Fig. 31. EDS determination of Ni/Cr vs 6 content.

338
TEXTURE IN STAINLESS-STEEL WELDS: AN ULTRASONIC STUDY

H. M. LEDBETTER, M. W. AUSTIN

Fracture and Deformation Division, National Bureau of Standards, Boulder,

Colorado 80303, USA

We studied texture effects in five AISI 316-stainless-steel welds. We measured

nine independent ultrasonic velocities along the weld's principal axes.. These

velocities reveal a strong texture different from the <100>-fiber-type usually

attributed to these materials. The texture is independent of delta-ferrite

content and must contain a strong <110> element related to the fee monocrystal

elastic constant C' = (C^-C^)/^. A postulated ideal ( 001 ) [110] plate-type

texture explains the range of measured longitudinal and transverse sound

velocities, but not their detailed values.

Key Words: Elastic constant; sound velocity; stainless steel; texture;

ultrasonics; weld

339

1. Introduction

Most metal processing produces texture, a nonrandom distribution of

crystallite orientations in a polycrystal. Well-known processes that produce

texture include: mechanical deformation, annealing (recrystal 1 ization)

,

solidification, electrodeposition, vapor deposition (epitaxy), phase

transformation, and welding. In each case one can identify a gradient

(mechanical, thermal, electrical, and so forth) that favors one crystal

orientation over another.

341
Texture induces anisotropic macroscopic physical and plastic properties.

Thus, usual polycrystal -monocrystal relationships for quasi- isotropic

(nontextured) aggregates provide only a rough guess for the macroscopic

material properties. A better guess must include information concerning the

orientation-distribution function of the crystallites.

In principle, any property affected by texture provides a possible tool

for probing texture. For example: electrical resistivity, thermal

expansivity, thermal conductivity, and so forth. But, these particular

properties, which one represents as second-rank tensors, can not work for the

cubic- symmetry case where they are simply direction-independent,

texture- independent scalars.

A texture-probing physical property useful for all crystal - symmetry

cases, including cubic, is the elastic-stiffness tensor. For the cubic case

it contains three independent components: usually C^, C^, and in

Voigt's contracted notation. Below, we consider the sound velocity, v

=

(C/p) 2

, where p denotes mass density. In general, sound velocities are roots

of the Christoffel equations:

det (c ijki

x

j

x

k

- = 0

where C..,, denotes the full fourth-rank elastic- stiffness tensor, x. denotes

ijkl j

components of the unit wave vector relative to cubic axes, and 6^ denotes the

Kronecker delta. Thus, in an anisotropic medium, v depends on direction

according to the elastic constants of the medium. Thus, in principle, one can

determine the symmetry and the orientation of an anisotropic elastic body by

measuring v in various directions.

In practice, one seldom uses the physical properties to study texture;

usually via x-ray diffraction one obtains orientations of particular

342
crystal lographic planes. Two references summarize the phenomenology of

texture and the x-ray-diff raction approach: Wasserman and Grewen [1] and

Barrett [2]. A single source dominates the description of mathematical

texture-analysis: Bunge [3], which contains 302 references.

The present study used elastic constants, or, equivalently, sound

velocities, to examine texture in 316-stainless-steel welds. These welds

contained up to ten volume percent of delta ferrite .

2. Materials

From a commercial source, we obtained welds in the form of 2.5-cm plates of

AISI-316 stainless steel. These were welded by a shielded-metal -arc process

under the following conditions: E316L electrode, 30 amperes, 24 volts, 15-20

cm/min travel speed, 93-212°C interpass temperature, 25-33 passes. Table I

gives chemical compositions. Table II gives some metallurgical properties.

Further material characterization together with some mechanical-property

measurements were given by Read et al . [4].

We prepared specimens for ultrasonic study by cutting, grinding, and

polishing using usual metal lographic procedures. These specimens were

rectangular prisms measuring approximately 2.5 x 2.5 x 0.3-0. 5 cm.

3. Sound-velocity measurements

Longitudinal-mode and transverse-mode sound velocities were determined by a

pulse-echo method described in detail previously [5]. Briefly, rectangular

prisms were prepared by grinding so that opposite faces were flat and parallel

within 5 ym. Quartz piezoelectric crystals with fundamental resonances

between 4 and 7 MHz were cemented with phenyl salicylate to the specimens. An

x-cut transducer was used for longitudinal waves and an ac-cut for transverse

343
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 = 2 £/ 1 (2)

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

oscilloscope, t was the time between adjacent echoes, the first and second

echoes usually being measured, and within these echoes the time between

leading cycles. Elastic constants were computed from the general relationship

C = pv

2

(3)

where p denotes mass density.

Extensional -mode sound velocities were determined by a Marx-osci 1 1 ator

method described elsewhere [6J. These measurements lead directly to Young's

modulus via the relationship

E = 4pf V (4)

where f denotes specimen resonance frequency and z denotes specimen length.

4. Results

Figure 1 shows measured sound velocities versus ferrite content. In this

figure, the first subscript indicates wave-propagation direction and the

second indicates polarization direction. The dashed lines in Fig. 1 indicate

the longitudinal and transverse sound velocities for nontextured 316 stainless

344
steel [7]. The co-ordinate system is: x^ perpendicular to welded plate, Xj

along weld direction, x

^ orthogonal to x^ and x^. For specimen 5, Table III

gives eight sound velocities; could not be determined accurately and is

not reported. (We assume v^^ = v^.) The set °'f: nine velocities v.. (i,j =

1,2,3) correspond to all possible waves propagated and polarized along the

principal axes x^x^, and x^. We measured no off-axis sound velocities. Also

for specimen 5, the Young's modulus results were E^ = 1.52 and =

11 2 1.20*10 N/m . Our specimens were too short along x^ to permit measurement of

E along the weld direction.

5. Discussion

Figure 1 shows that the sound velocities and, therefore, the elastic constants

of these stainless-steel welds do not vary strongly with composition. This

suggests that the elastic constants of delta ferrite do not differ

significantly from those of austenite. This is hardly surprising because the

atomic volumes of the two phases differ by only one or two percent. Also,

Ledbetter and Reed [8] showed for Fe-Ni alloys that the alpha-ferrite and

austenite elastic constants do not differ dramatically.

For all cases, the longitudinal sound velocity v^ falls below the

isotropic (nontextured) value, 0.569. This suggests the existence of texture

in all the specimens. It is consistent with the <100> fiber texture often

suggested for these welds [9], Ledbetter [10] calculated that a pure <100>

fiber texture requires v^ = 0.510.

One expects that cubic- symmetry crystallites in welds occur either with

random orientations or with a single preferred axis parallel to the principal

heat-flow direction, denoted x^. In the x^-x^ plane, crystallites are

expected to orient themselves randomly because no preferred heat-flow

direction occurs in that plane. Thus, an effective fiber texture occurs with

345
being the fiber axis. Macroscopical ly , such a material exhibits

transverse-isotropic symmetry, with five independent elastic constants. This

is equivalent to hexagonal symmetry. In Voigt's contracted C.. notation the

symmetrical elastic- stiffness matrix is:

'11 '12

'11

,H Q

'13 u

,H'13 0

0 0

0 0

'33

'44

0 0

0 0

C Hl o 44

u

M

'66

(5)

H 1 H H where - C^)- Thus, the theoretical problem is to express the

|_| five hexagonal -symmetry elastic constants, C.., in terms of the three cubic1 sJ symmetry elastic constants: C^, C^, C^.

Present theory does not permit calculation of the elastic constants of

textured polycrystal 1 ine aggregates of cubic-symmetry crystallites. This is a

special and more difficult case of determining the average elastic constants

of a random polycrystal 1 ine aggregate. For this latter problem, Landau and

Lifshitz [11] remind us that "there is ... no general relationship between the

moduli of elasticity of a polycrystal and those of a single crystal of the

same substance." The fiber- texture case involves a two-dimensional, rather

than a three-dimensional, average.

However, elastic-constant bounds can be calculated for textured

aggregates. In this study we calculate the Voigt (constant strain) and Reuss

346
(constant stress) bounds and use Hill's suggestion of an arithmetic average to

estimate the elastic constants. This approach turns out to be quite

reasonable because the Voigt and Reuss bounds do not differ dramatically (a

maximum of 37 percent). And for the important <100>-texture case, the bounds

H H H coincide for three elastic constants: C^» C

33 , and C 44< Thus, we know these

three constants unambiguously. Kroner and Wawra [12] explain the systematics

of this bounds coincidence, and they give the following expressions for the

C

H

.: 1J

where

C

H

L 11

= Cj + 3 BC^ = 4< 3C

11 + C12 + 2C44 )

c

H l

12

= + BC

3

= 4< c u +

3C^

2

- 2C 44 )

C Hl

13

= C

2

- 4bC

3

= C

12

c

H

l

33

= Cj = 8 bC

3

= C 11

C Hu 44

= \(c

1

- c

2

) - 4B C 3

= C 44

H

C

66

=

!/

r H pH x 2^ii - ^ 12

1

= J(c 11

‘ C 12 + 2C 44

5C

1 = 3C

11 + 2C

12 + 4C44

5C

2

= C

11 + 4C

12 + 2C 44

C

3

= C

11

" C

12

' 2C

44

( 6 )

( 7 )

where

B = 1/20 for <100> fiber texture.

Apparently, relationships for the S.., which give the Reuss bounds by

matrix inversion, do not appear in the literature. These are:

347
H

11 b

(6s h

+ 2S

12

+ W

H12 S

(2S

11 + 6Sj2 • S^)

H13

= S

12 ( 8 )

S = S5

11

S = S ^44 ^44

.H S

66

2 ^ S

11

" S 12^ = S

11

' S

12

’ S 44/2

-1

In deriving these relationships, one must remember that the inverse C^i of a

fourth-rank elasticity tensor relates to the inverse C"J of its associated 6 x

6 matrix by a factor multiplied times . This factor is 4 if both a and B

exceed 3, is 2 if either a and B exceeds 3, and is 1 if neither a nor B

exceeds 3.

The most dramatic result of our measurements, as shown in Fig. 1, is the

splitting of the isotropic v^ = = 0.307 shear velocity into two

velocities v^ = 0.381 and v^ = 0.228, on average. One expects the

shear-wave birefringence to provide a more sensitive probe of texture than the

smaller shift in the longitudinal velocity. This arises because, even for

highly anisotropic cubic crystals, v

£ depends only slightly on direction [13].

However, this particular v^ splitting is inconsistent with a <100> fiber

texture for which Ledbetter [10] showed that the two velocities should be

0.384 and 0.289. That is, v^ = 0.228 is much too low for a <100> fiber

texture.

Departure of measurement from a <100> fiber texture is shown

schematically in Fig. 2. At the top. Fig. 2 gives the bounds of Voigt (upper)

and Reuss (lower). Any type of texture in any amount must produce velocities

348
within these bounds. Our measurements meet this criterion. For the shear

mode, we found velocities very close to both the upper and lower bounds, which

correspond to and to C' = (C^ - respectively. Thus, the fee

monocrystalline and C' must be two of the principal elastic constants of

the textured aggregate. Again, the often-suggested <10Q> fiber texture fails

as a possible texture because its elastic constants contain but not C'.

If we assume another texture element, in addition <100> in the weld direction,

the situation worsens, as Fig. 2 shows. The situation improves if we assume

<110> in the weld direction; this amounts to a (001)[110] texture. Figure 2

shows that this ideal texture explains the range of shear velocities, but not

the detailed measurements without some perturbations. Possible perturbations

include partial texture, a second texture, and deviations from a (001) plane

or a [110] direction (or both).

6. Summary

The present study produced the following conclusions concerning the elastic

constants and sound velocities of 316-stainless-steel welds:

1. Up to at least ten-percent delta-ferrite content, the elastic

constants are independent of ferrite content.

2. All specimens exhibit texture and significant departures from

quasi-isotropic physical -property values. The texture is

independent of ferrite content.

3. For these materials, the usual model, <110> fiber texture, fails to

explain the measurements.

349
4. Calculated results for several other textures show that the elastic

constant C' = (C^-C^)/2 or a [110] texture element, enters

essentially into the weld's texture.

5. A postulated texture -- (001)[110] -- covers the range of eight

measured sound velocities, but not the details of their

distribution. Some other, presently unknown, element of texture

must also occur.

Acknowledgment

This study arose from research sponsored by the U.S. DoE Office of Fusion

Energy.

References

[1] G. Wasserman and J. Grewen, "Texturen metallischer Werkstoffe"

(Springer- Verlag, Berlin, 1962).

[2] C. S. Barrett, "Structure of Metals" (McGraw-Hill, New York, 1952).

[3] H.-J. Bunge, "Texture Analysis in Materials Science" (Butterworths

,

London, 1982).

[4]

, D. T. Read, H. J. McHenry, P. A. Steinmeyer, and R. D. Thomas,

Weld Res. J. Suppl . 59 (1980) 104s.

L5]. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl . Phys.

51 (1980) 305.

[6] H. M. Ledbetter, Cryogenics 20 (1980) 637.

[7] H. M. Ledbetter, Metal Sci . 14 (1980) 595.

L8] H. M. Ledbetter and R. P. Reed, J. Phys. Chem. Ref. Data 2 (1973)

531.

[9] B. R. Dewey, L. Adler, R. T. King, and K. V. Cook, Exper. Mech. 17

(1977) 2.

[10] H. M. Ledbetter, in "Review of Progress in Quantitative

Nondestructive Evaluation, Volume 1" (Plenum, New York, 1982)

p. 619.

350
[11] L. D. Landau and E. M. Lifshitz, "Theory of Elasticity" (Pergamon,

London, 1959) p. 40.

[12] E. Kroner and H. H. Wawra, Philos. Mag . A38 (1978) 433.

[13] H. M. Ledbetter and R. L. Moment, Acta Metal 1

.

24 (1976) 891.

351
Table I. Chemical compositions, weight percent, balance Fe

Specimen Cr Ni Mo Mn C N Nb Si S P

1 18.47 13.97 2.25 1.72 0.031 0.042 0.14 0.58 0.011 0.033

2 18.64 12.97 2.29 1.72 0.030 0.046 0.14 0.56 0.012 0.034

3 19.20 12.01 2.32 1.70 0.034 0.042 0.15 0.57 0.009 0.034

4 19.66 11.06 2.18 1.76 0.028 0.067 0.15 0.59 0.008 0.029

5 18.69 12.85 2.25 1.75 0.059 0.044 0.15 0.56 0.009 0.033

Table II. Metallurgical properties

Specimen

Ferrite

(vol ume percent)

Ferrite

Number

Mass Density

(g/cm3

)

1 0.13 0.12 7.932

2 4.1 4.5 7.934

3 8.5 9.2 7.921

4 10.1 11.0 7.910

5 1.2 1.3 7.930

352
Table III. Sound velocities in specimen 5 in units of cm/ys

11

= 0.555 v 22= 0.568 v 33

= 0.538

12

= 0.222 V

21

= 0.260 V 31

= 0.363

13

= 0.381 ( y 23= 0.387) v 32= 0.387

353
Figure Captions

1. Longitudinal and transverse sound velocities versus delta-ferrite

content for 316-stainless-steel welds. Dashed lines indicate

isotropic (nontextured) velocities: upper = longitudinal, lower = transverse.

2. Longitudinal and transverse sound velocities for 316-stainless-steel

welds. Small numerals indicate n-fold degeneracy.

354
(st/

/mo)

A1I00T3A

ONHOS

1. Longitudinal and transverse sound velocities versus delta-ferrite

content for 316-stainless-steel welds. Dashed lines indicate

isotropic (nontextured) velocities: upper = longitudinal, lower = transverse.

355
SOUND VELOCITY (cm//is)

0.2 0.3 0.4 0.5 0.6 0.7

r

1 1 1 1 1'

Bounds

3

Isotropic

Measurement

2

tool] Fiber

T exture

( 001 ) [100]

T exture

( 00

1

) [ 11

0

]

Texture

Longitudinal and transverse sound velocities for 316-stainless-steel

welds. Small numerals indicate n-fold degeneracy.
HOMMETALUCS

357

NONMETALL ICS PROGRAM

LEADERS: M. B. Kasen, H. M. Ledbetter, NBS

STAFF: M. W. Austin, R. D. Kriz, R. E. Schramm, NBS

CONTRACTS: S. K. Datta, University of Colorado

OBJECTIVES:

(1) To measure, model, and understand the properties of composites

that enter essentially into the conception, design, construction, operation, and performance of magnetic-fusion-energy

equipment.

(2) To cooperate with industry to establish and refine

cryogenic-grade specifications for insulating laminates and

coding systems for laminates.

RESEARCH PAST YEAR (1982):

(1) Kriz, Ledbetter. Ultrasonic measurement of the complete

nine-component elastic-constant tensor of a woven

glass-fiber/epoxy composite.

(2) Kriz. Calculated stress distributions in a woven composite by

a finite-element model.

(3) Datta, Ledbetter, Kriz. Predicted elastic constants of

transversely isotropic composites containing anisotropic

fibers.

(4) Ledbetter, Kriz. Elastic-wave surfaces in fiber- reinforced

composites.

(5) Datta, Ledbetter, Austin, Kriz. Longitudinal wave propagation

in an elastic composite thin rod.

(6) Kasen, Schramm. Screening, down to 76 K, of samples of G-10CR

produced by two new manufacturers of this product.

(7) Kasen. Continuation of a program on the standardization of

cryogenic grades of insulating and structural laminates in

cooperation with industry.

RESEARCH THIS YEAR (1983):

(1) Kriz, Ledbetter, Datta. Modeling of laminated composite

elastic properties.

(2) Ledbetter, Austin. Anisotropic thermal expansion of

glass-fiber/epoxy composites: measurements and modeling.

(3) Ledbetter, Datta. Anisotropic thermal expansion of uniaxial,

fiber- reinforced composites: measurements and modeling.

(4) Kasen. Development of improved test methods for determining

the effect of cryogenic temperatures on composite material

performance.

359

Variability in Mechanical Performance of G-10CR

Cryogenic-Grade Insulating Laminates*

M. B. Kasen and R. E. Schramm

Fracture and Deformation Division

National Bureau of Standards

Boulder, CO 80303

ABSTRACT

An assessment is made of the variability in cryogenic mechanical

performance of insulating laminates produced by five manufactureres to a common component and procedure specification. Results at 295 K and 76 K

indicate that the specification is adequate to ensure the desired degree of

product uniformity.

* Published in Cryogenics, Vol . 23, No. 5, 273-280 (1983).

361

Five United States firms have produced a controlled grade of glass-fabric

reinforced epoxy high pressure industrial laminates designated G-10CR. This

laminate grade was originally intended to provide designers of superconducting

magnets for magnetic fusion energy research with commercial sources of an insulating laminate having minimum performance variability at 4 K. Laminate

production is controlled by component and manufacturing specifications

developed cooperatively by industry and the National Bureau of Standards.

The basic mechanical, thermal, and electrical performance of this grade has

2

been investigated; however, little information has been available on the

variability in performance of laminates produced to the specification by

different companies. NBS therefore undertook to assess this variability by

comparing the effect of cryogenic temperatures on tensile, compressive, and

interlaminar shear properties of G- 1 OCR products produced by the five

companies. Only the warp direction of reinforcement in the plane of the

laminate was considered, and testing was confined to 295 K and 76 K.

2 The results, obtained by procedures described elsewhere, are summarized

in Table 1. This Table reports averages, standard deviations, and coefficients of variance for the nine to eleven specimens tested in assessing each

parameter. Lower tolerance limits on the properties are also reported, based

on confidence limits of 99% and a confidence coefficient of 0.95. Assuming

that the distribution of test results is normal, and that the materials

evaluated are representative of commercial products produced by each company,

there is a 95% probability that 99% of the production will produce values

higher than the lower tolerance limits. The assumption of a normal distribution of strength properties may be questioned-- in particular, at cryogenic

temperatures where the matrix becomes embrittled. However, there were insufficient data available to justify assuming a different distribution, and the

error introduced by this assumption is probably small. In Table 2, the

Contribution of NBS, not subject to copyright in the U.S.A.

363
mechanical property data are given separately for the G-lOCr products of each

manafacturer.

Data on the coefficient of variation and on lower tolerance limits for

tensile strain at failure are not reported because failure of the strain gages

at stresses below maximum required that this parameter be estimated from a strip chart recording. This significantly reduced the accuracy of the

measurements, and precluded determination of a meaningful variability.

When applied to a single production lot of G-10CR, our test procedures

produced coefficients of variation of about four percent for tensile strength,

Young's modulus, and shear strength, and about nine percent for compressive

strength, Poisson's ratio, and failure strain. This suggests that the co- efficients of variation in the present study primarily reflect uncertainties

in the test procedure and intrinsic material variability. It is therefore

concluded that differences in room or cryogenic mechanical performance of

products produced to the G- 1 OCR specification by different companies are probably not statistically significant. The specification therefore appears

adequate to ensure the desired product uniformity. The coefficient of

variation is higher at 76 K than at 295 K for most properties. This is an expected consequence of matrix embrittlement, which severely limits the

operative relaxation processes, making the strength properties more sensitive

to the presence of small flaws. Residual stresses combined with the

possibility of microcracking also increase the variability in the elastic

properties at cryogenic temperatures as compared to room temperature.

It should be emphasized that G-10CR products are highly anisotropic, the

values for the warp direction of reinforcement being the highest available in

the plane of the laminate. Performance in the fill direction will be substan2

tially lower , as will the performance for stresses applied in orientations

intermediate between the warp and fill directions. Furthermore, this study

was limited to evaluation of cryogenic performance at 76 K, whereas the

functional use of the materials is at 4 K. Testing at 4 K would have yielded

2

somewhat higher strengths and moduli , but would also be expected to increase

the coefficient of variation.

There is reason to believe that the average and lower tolerance limits

for the compression and interlaminar shear properties in Table 1 may be

somewhat lower than those representati ve of bulk material properties. Recent

data have suggested that the compression specimen used in this study may be

364
sensitive to edge effects, causing premature failure. There is also reason to

believe that the guillotine method used to evaluate interlaminar shear in this

study yields values that are somewhat below the true bulk shear values. It is

therefore probable that these values are conservative.

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

of Energy.

Terminology

Oj = Ultimate tensile strength

c?£ = Ultimate compressive strength

Ej = Young's modulus in tension

v = Poisson's ratio

£j = Ultimate failure strain in tension

a i is

= Ultimate interlaminar shear strength

References

1. Kasen, M. B., in Nonmetal lie Materials and Composites at Low

Temperatures-2, G. Hartwig and D. Evans, eds., Plenum Press, NY (1982),

327.

2. Kasen, M. B., MacDonald, G. R., Beekman, D. H., and Schramm, R. E., Adv

Cryo Eng , 26 (1980) 235.

365
Table 1. Performance of G-10CR cryogenic-grade insulating laminates

Property Average

Standard

Deviation

Coefficient of

Variation

Lower Tolerance

Limit*

°T

431 MPa

295 K

27.0 MPa 6.26% 306 MPa

°C

390 MPa 30.8 MPa 7.89% 254 MPa

e t

27.8 GPa 1.29 GPa 4.64% 22.0 GPa

V 0.159 0.014 8.67% 0.098

e T

1.97% 0.35% — --

°ILS

48.1 MPa 2.48 MPa 5.17% 33.7 MPa

°T

817 MPa

76 K

71.4 MPa 8.74% 500 MPa

°C

703 MPa 52.1 MPa 7.41% 472 MPa

e t

32.1 GPa 2.72 GPa 8.47% 20.1 GPa

V 0.208 0.023 11.2% 0.106

e T

3.05% 0.395% --- —

°ILS

68.7 MPa 5.99 MPa 8.72% 43.0 MPa

^Confidence limits 99%

Confidence coefficient 0.95

366
Ultimate

Ultimate

Compressive

Interlaminar

Shear

Interlaminar

Shear

Temp,

K

Young's

Modulus

Poisson's

Ultimate

Tensile

Strength

Tensile

Strength

Strength

(Guillotine)

Strength

(Short

Beam)

GPa

Msi

Ratio

MPa

Ksi

Strain,

%

MPa

Ksi

MPa

Ksi

MPa

Ksi

00 r-* O LO 00 00

00 LO

os os

os lo

00 OS

as lo

LO oo

LO LO co oo

oo

LO LO O LO

oo oo

HCO C\J as r^. lo

•K * —• oo O CM

LO LO A A

OO O OO LO

* * O<3- O CM LO O O

cm

i—< O'* lo

oo O o

OOOi oo oo o cn o lo COOOi

•-* CO

LO

* *

<3- O

O LO

00 LO

* *

<3- CO

00

LO

os r>.

<3- oo oo

<3- <3*

r-H CJY LO LO LO LO

ro a.

as cm

os o

O'* 00

<3- <3- LO LO o cr* *—* LO o —

•

LO

LO *—

•

as os

*oo

s- Q_

«-h CO lo CO o CO CO LO N

as <3- cm co

CM «—• 00 IM ^ O IM

Q_

s.

<T3

C\J N CO LO

*3* CO CO LO

co co lo

LO O OS LO

co co

CO

cr. os

lo MN o CO O LO co CO «3- LO LO

co co

CO CO

00 LO

r-H CM

<3- 00 CM CM

00 LO

*—• CM

CO os

CO CM HD 21

00 o

—i CM

osCM CM

OS LO

CM CM O «-H LO LO

LO 00

CM CM

LO LO O O

OS «—I LO

CM OS LO ^ 00 00

.—• cm i—

*

•—•CM 1^ 00 00 00 o co *—

•

•—•OS LO CM

co

r^. as

us is i\ co ^ 00 CO

OS < r—

t

*—• LO co co

<3- ^ i—

.

O O co «—

i

co •—•co

r—* r—• CM CM

<3- o LO LO rH CM

•—• •—» CM CM ^ •—• 00 LO Nl/1 r—i •—* •—• •—• CM CM

CM LO LO LO LO O •—

•

•—• r—i CM CM o o o o o o o o o o o o O O O O O O O O

•—• o oo

OS 1—I LO

co co os o O cmcm co CM

co co

r-H OS CO CM

CM CM 00 O

LON r^» CM O 00 —• «-H co Cj- ^ ^ LO co co lo CO LO LO

^ co co 00 CM O CM (ON OS oo O LO CO LO O co

co o cm

CM CM CO co

r^. as co lo

CM CM CO CO

LO LO CM OS

CM CM CO CM

as os co *3

- CM CM CO CO

00 LO LO LO

CM CM CO CO

LO LO

OS N*. CM

LO LO

osCM

LO LO

OS N'- CM

LO LO

os

CM

K

LO LO

osCM

367

Broke

in

tension,

not

shear

* EFFECT OF CRYOGENIC TEMPERATURES ON THE MECHANICAL PERFORMANCE OF

CLASS-FABRIC-REINFORCED EPOXY AND POLYIMIDE MATRIX LAMINATES

M. B. Kasen, R. E. Schramm, and R. D. Kriz

Fracture and Deformation Division, National Bureau of Standards, Boulder, Colorado, USA

Radiation-resistant laminates are required for insulators and

structural supports in the superconducting magnets of magnetic

fusion energy systems. Glass-reinforced laminates fabricated

with a polyimide matrix have a much higher tolerance to neutron

and gamma radiation at 4 K than do laminates fabricated with an epoxy matrix. However, tests indicate that the mechanical

performance of polyimide-matrix laminates is inferior to that of

the epoxy type, whereas their elastic performance is superior.

Fractographic studies show that the performance difference is due to a lower integrity of the polyimide-glass interface.

* Published, Proceedings of the International

Japan, 11-14 May 1982, Editors K. Tachikawa

Cryogenic Materials Conference, Kobe,

, A. Clark (Butterworths , UK, 1983).

369

INTRODUCTION

Conventional epoxy-matrix glass-reinforced insulating laminates have been used for

many years in superconducting magnet construction, but recent studies have shown

that such laminates degrade rapidly when exposed to 100 kGy (1 Gray «= 1 J/kg

100 rads) of gamma radiation at 5 K [1]. Poly imide-matrix laminates degrade much

less rapidly [2]; however, such laminates have disadvantages in basic performance

and cost. This paper describes a program comparing the effect of cryogenic temperatures on the mechanical and elastic performance of the two laminate types. Our

materials and those for the radiation tests were from the same manufacturer.

MATERIALS AND PROCEDURES

This study compares performance of a standard bisphenol A epoxy-matrix cryogenic

grade of high-pressure industrial laminate, G-10CR [3,4], with that of a laminate of similar construction utilizing a polyimide matrix, designated PG-10CR in this paper.

Tensile, compressive, and interlaminar shear strength, elastic modulus and Poisson's

ratio, and ternsile ultimate strain were determined for each at 295 K, 76 K, and 4 K using test procedures previously described [3]. Data from two specimens of PG-10CR

are compared with data from three to six specimens of G-10CR f»r each test condition

at each temperature. Both warp and fill direction properties ^f the laminates were

studied.

A scanning electron microscope (SEM) f ractographic examination was performed on specimens of each laminate type fractured in tension, compression, and shear at each

temperature to associate failure modes with property variability.

RESULTS

Mechanical Properties

Figure 1 illustrates the range encompassed by tests in the warp direction for the

two laminate types. Tensile and compressive strength of the polyimide laminate were

371
found to be 20-402 lower than that of the epoxy laminate at cryogenic temperatures.

Interlaminar shear strength of the polyimide laminate wag lower by the same amount

at all temperatures as was the strain-to-failure in tension. Young's modulus of the polyimide laminate was 10-202 higher than that of the epoxy laminate at all temperatures.

Fractography

Fractographic evidence suggested that the lower strength of the polyimide laminate

at cryogenic temperatures was primarily due to low integrity of the polylmlde-glass

interface. As illustrated by Fig. 2, the polyimide laminate shattered during

tensile failure at cryogenic temperatures, leaving the fibers near the fracture

almost entirely devoid of resin and causing severe delamination well away from the fracture. By contrast, Fig. 3 illustrates that the epoxy laminate failed much more

locally under the same conditions, with fiber bundles retaining a large amount of resin. The higher magnification views in these figures clearly show that the polyimide laminate failed cleanly at the glass surface, and the epoxy laminate

failed through the resin at some distance from the glass surface. This difference

in failure mode is clearly temperature dependent, since relatively little difference

in tensile failure mode was found between the two laminate types at room temperature

.

The lower interlaminar shear strength of the polyimide laminate is also attributable

to inability of the interface to sustain appreciable shear stresses. The failure

surface was again characterized by clean glass fibers and absence of the resin

matrix. In contrast, the failure surface of the epoxy laminate showed a high degree

of resin retention. Fractographic examination showed that the high interlaminar

shear strength of the epoxy laminate at cryogenic temperatures was caused, in part,

by tensile failure of some glass fibers, evidencing the ability of the glass-epoxy

interface to transfer shear stresses into the fibers efficiently.

CONCLUSION

The inferior mechanical performance of the glass-polyimide laminate, as compared

with that of the glass-epoxy laminate, is attributable to poor integrity of the glass-polyimide interface. This deficiency is exacerbated by cooling to cryogenic

temperatures and is particularly damaging to the shear properties of the laminate.

LIST OF SYMBOLS

T(K) = temperature, kelvins

E = Young's modulus

v = Poisson's ratio

c = tensile ultimate strain

TU

= tensile ultimate stress

o^y = compressive ultimate stress

= interlaminar shear ultimate stress

ACKNOWLEDGMENTS

This work was supported by the Office of Fusion Energy, U.S. Department of Energy. We thank Mr. J. R. Benzinger of the Spaulding Fibre Company for providing the test materials.

372
references

Coltmart , R.R., Klabunde, C.E., Kernohan, R.H., and Long, C.J. Radiation Effects

on Organic Insulators for Superconducting Magnet s, 0RNL/TM-707 7 , Oak Ridge

National Laboratory, Oak Ridge, Tennessee, 1979.

Coltman, R.R., Klabunde, C.E., and Long, C.J. The Effect of a 100 MGy

(IQ 10 rads) Gamma-Ray Dose at 5 K on the Strength of Polylmlde Insulators , Report SSD No. 81-10, Oak Ridge National Laboratory, Oak Ridge, Tennessee,

1981.

Kasen, M.B. , MacDonald, G.R., Beekman, D.H., and Schramm, R.E. "Mechanical,

electrical and thermal characterization of G-10CR and G-11CR glass-cloth/epoxy

laminates between room temperature and 4 K," A.F. Clark and R.P. Reed (eds.),

in Adv ances in Cryogenic Engineering (Materials) , vol. 26, pp. 235-244, Plenum

Publishing Corporation, New York, 1980.

Benzinger, J.R. "Manufacturing capabilities of CR-grade laminates," A.F. Clark

and R.P. Reed (eds.), in Advances in Cryogenic Engineering (Materials) , vol. 26, pp. 252-258, Plenum Publishing Corporation, New York, 1980.

1000r

&

z

800

600

400

200

- a

Termite Strength

(werp)

a

4 K 76 K 295 K

1000

nU 800

CO A _

600

3

, O

400

200

Compressive Strength

(werp)

0.3 r

4 K 76 K 295 K

& 0.2

0.1

Poteeon’e Ratio (warp) m

4 K 76 K 295 K

42 f Young's Modulus

(warp)

38 -

S. 34

o

I1J 30

26

22

4 K 76 K 295 K

85p Shear Strength

(warp)

75-

« 65

Q»

3 55

b

45

35

4 K 76 K 295 K

G-10CR

PG-10CR

5 r

- 3 ?

w

Strain to Failure (warp)

0

4 K 76 K 295 K

Fig. 1. Comparative range of mechanical properties for epoxy matrix

polyimide matrix (PG-10CR) laminates.

(G-10CR) and

373
a Is

Fig. 2.4-K tensile failure, glass-polyitnide

.

Fig. 3. 4-K tensile failure, glass-epoxy.

374
ELASTIC STIFFNESS AND LOCAL STRESSES IN WOVEN-FABRIC

COMPOSITES AT 77 K

Ronald D. Kriz

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

375

Woven-fabric composites are commonly used in superconducting magnets

and in containment of cryogenic liquids. Here, the mechanical response

of a plain-weave laminated composite at cryogenic temperatures is studied

by predicting the load-deformation response of a fundamental "unit cell".

The unit cell is the smallest repeating geometrical structural unit such

that 90-degree rotations about the axis normal to the fabric plane yield

four identical patterns. The model assumes that generalized plane strain

exists on all four faces of the unit cell. Hence the elastic response of

the unit cell can be predicted by studying the three-dimensional loaddeformation response within a single plane parallel to the load direction

as shown in Fig. 1. Stiffness and stresses within this plane are

calculated for worst-case load conditions prior to fracture. A finiteelement model is used to predict unit-cell stresses and stiffnesses.

Several load conditions of tension, shear, and compression are

possible. Preliminary results given here focus on tensile loads, N ,

shown in Fig. 1. The displacements within the x-z plane are linear

functions of x, y, and z coordinates and out-of-plane strain, e^:

377
u a^ + a£X + a^z

v ( 1 ) = a 4

+ a 5

x + a 6z + ye

y

w = a^ + agX + agZ

where the a. terms are evaluated in terms of nodal displacements. The

in-plane and out-of-plane strains, e and e , are calculated from thinx y

laminate-plate theory assuming zero stresses through the thickness:

E x

* W<AxxVA

xy>’

ey

= AyV<AxxVA

xy> < 2 >

where A , A , and A are laminate midplane stiffnesses. Here a xx yy xy

laminate load (N

x =1751N/cm) yields ex

= 0.01095 and = 0.001741 with

elastic properties taken from Table 1. A worst-case thermal load of

AT = - 318 K is chosen at 77 K assuming a strain-free state at 395 K

(cure temperature)

.

With these load conditions we study the elastic mechanical response

and internal stress states of the unit cell. In Fig. 2 the influence of

warp-fiber orientation, 0, on the elastic response of the unit cell is

shown. As expected, the elastic modulus increases with decreasing angle.

Regions of highest stress concentration prior to fracture are observed

along lines a-b and c-d of Fig. 1. These stress distributions are

plotted in Figs. 3 and 4. As expected, larger compressive stresses are

observed along line a-b of Fig. 3 with increasing warp angle. The

influence of the thermal load increases the compressive stress for the

45° case. Stresses normal to the warp interface increase with increasing

warp angle; the largest effect occurs in the matrix region. The thermal

378
load increases the 45° normal stresses in the matrix region and decreases

these stresses in the fill region. Thus, the matrix region is most

affected by thermal loads and warp-fiber orientations.

We continue parametric studies of the load-deformation response of

the unit cell for other load cases, including the redistribution of

stresses caused by the presence of cracks in the fill fiber bundles.

This model permits us to evaluate other elastic properties such as shear

modulus in the x-z plane. We plan to report these results later.

379
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Volume

fraction

(Vf)

of

glass

fibers

Subscripts :

(I)

indicates

fiber

direction

and

(t)

indicates

transverse

plane
X

381

Thin

slice

of

unit

cell
5 Msi

4

3

Fig. 2 Unit cell modulus in x direction (warp)

382
-25 ksi

0.2 0.3

Z / T

0.4

0

Fig. 3 z stresses along line a-b in Fig. 1.

383
BdW ‘ JQ

384

Fig.

4

Stresses

normal

to

warp-fill

interface

along

line

c-d

in

Fig.

1.
Influence of Ply Cracks on Fracture Strength

of Graphite/Epoxy Laminates at 76 K

Ronald D. Kriz*

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

Abstract

Quasi-isotropic laminates ([0/90/±45] ; [0/±45/90] ) were fabricated from

graphite/epoxy and quasi-statically loaded s in tension at 76 K until fracture

occurred. Fibers in 0° plies carry the largest portion of the tensile load;

the weaker 90° and 45° plies crack at loads much lower than fracture strength.

The effect of ply cracks on fracture of load-bearing 0° plies was examined to

understand how defects affect laminate strength. A generalized plane-strain

finite-element model was used to predict stress gradients in the 0° ply near

the crack tip. Variations in residual stress caused by changes in temperature

and absorbed moisture were included in the analysis. The experiments

indicated that absorbed moisture significantly alters the fracture strength

and fracture surface of a dehydrated [0/90±45] laminate tested at 76 K. The

0

C plies of dehydrated laminates fractured alofig several 90° ply cracks. When

moisture saturated a [0/90/±45] laminate, a single 90° ply crack dominated

the fracture of the 0° ply and the laminate fracture strength decreased 8

percent. Analysis of residual stresses indicated a higher residual stress

state existed near the 90° ply crack when moisture was absorbed.

Submitted for publication in ASTM special technical publication from ASTM

Symposium on Effects of Defects in Composite Materials, San Francisco,

California, December 1982.

Key Words: finite-element; fracture; graphite/epoxy laminates; low

temperatures; ply cracks; residual stress

* NRC-NBS Postdoctoral Research Associate 1980-1982

385

INTRODUCTION

When loaded in tension, graphite/epoxy laminates accumulate damage in the

form of ply cracks, delaminations, fiber breaks,' fiber-matrix debonding, and

matrix cracks. Prior to fracture, 0° plies carry a larger portion of the

tensile load than the weaker 90° and 45° plies. Laminate fracture occurs when

fibers fracture within the load-bearing 0° plies.

Experimental evidence of graphite fiber breaks occurring in 0° plies prior

to fracture was observed by Kriz [1] at room temperature. An enlargement of

these fiber breaks is shown in Fig. 1. The fiber breaks observed in Fig. 1

are not influenced by the stress concentration of a 90° ply crack. The

scatter in graphite fiber breaks is similar to those modeled by Rosen [2] for

unidirectional glass/epoxy. However, Hartwig [3] observed an accumulation of

graphite fiber breaks near 90° and 45° ply cracks at 76 K. Hence the presence

of large residual stresses near a ply crack at 76 K can influence the fracture

of 0° pi ies

.

In this study, the influence of residual stress on 0° ply fracture was

investigated for two quasi- isotropic laminates: [0/90/±45]

s and [0/±45/90] .

Residual stresses were altered, prior to fracture at 76 K, by dehydrating and

saturating laminates with absorbed moisture. A statistical analysis of

laminate fracture populations was designed to model the weakening of

load-bearing 0° plies. A generalized plane-strain finite-element model was

used to predict variations in 0° ply stresses caused by mechanical and

thermal -moisture loads.

387
ANALYSIS

Stresses within each layer were predicted for a symmetrical laminate

loaded in tension. Material response was assumed to be linear and elastic at

76 K. Hence, thermal and moisture loads were superposed with mechanical

loads. The applied loads were assumed to be steady and uniform across the

laminate. A typical symmetric laminate loaded in tension is shown in Fig. 2,

where L>B>>T. Each layer was assumed homogeneous with transversely isotropic

elastic properties. Room- temperature and 76- K elastic properties used in the

analysis are listed in Table 1. Coefficients of thermal and moisture expansion

(listed in table 1) were modeled as the only contributions to residual stress.

Changes in residual stress caused by variations in elastic properties were

neglected.

Generalized Plane Strain

Stresses near a ply crack tip (x=y=0, z/T=0.75) were predicted within the

x-z plane far from the laminate edges (y=±B). The constraining out-of-plane

strain, e , was modeled as the generalized strain normal to the x-z plane at

y=0. The strain in the load direction, e , was assumed to be uniform. Using

X classic laminate-plate theory (see Jones [5]), the strains e and e can be

x y

calculated in terms of the known laminate loads, N , with N = 0.

x y

e = A N /(A A -A 2

) , e = -A N /(A A -A 2

) x yy x

' xx yy xy' * y xy x ' xx yy xy' ( 1 )

where A , A and A are laminate midplane stiffnesses. Since generalized

xx* yy xy

K 3

plane-strain was assumed normal to the x-z plane, all six components of stress

must be independent of the y-axis. Hence the equations of equilibrium reduce

to

3a

x 3x

+

3 T xz

3Z

3t 3t 3t 3a

= _J<I + _J

3x 3z 3x 3Z

= 0. ( 2 )

388
The displacement relations are derived by Talug [6] for the generalized

plane-strain problem assuming laminate symmetry:

u(x,y,z) = U(x,z)

v(x,y,z) = e^y + V( x , z) (3)

w(x,y ,z) = W(x,z)

where U, V, and W are unknown displacement functions. Talug [6] derived the

elliptic equations of equilibrium in terms of these unknown functions and

solved for displacements near ply cracks using finite difference techniques.

In this study thermal and moisture loads were included in the analysis and a

finite-element model was used to solve for displacements near ply cracks.

Finite-Element Model

Constant-strain elements with three nodes were chosen. Hence, the

required displacement fields (3) are satisfied for each element by the

1 inear relations

u = a

l + a

2 x + a 3

z

v = a 4 + a

5 x + a 6

z + e y

y (4)

z = a

7 + a g

x + agZ

where the a. terms are evaluated in terms of nodal displacements. Elemental

stiffness and load matrices for mechanical and thermal -moisture loads were

derived by Renieri [7]. A finite-element computer program was written using

these elements. Because of laminate symmetry only the cross-hatched quadrant

shown in Fig. 2 was modeled. The grid of elements shown in Fig. 3 was used to

model stress concentrations within the 0° ply near the 90° ply crack. The

dimension of the smallest element at the crack tip was less than one

graphite-fiber diameter (see Region A of Fig. 3).

389
Traction-free boundary conditions along z=T are approximated by

prescribing statically equivalent zero nodal forces. Laminate symmetry along

z=0 requires displacements w=0. Similarly, displacements u=0 along x=0 are

required, except for the traction-free surface 0.5<z/T<0.75 that requires

statically equivalent zero-nodal forces. In the x-direction along the surface

x=L a uniform axial strain e is modeled by prescribing a uniform nodal

A displacement Uq = e x L.

To check the accuracy of the grid shown in Fig. 3 finite-element stresses

predicted with no ply crack were compared with thin- laminate- pi ate theory.

This comparison is shown in Table 2 when a grid aspect ratio of L/T =6 is

prescribed. Using the grid shown in Fig. 3, finite-element predictions of

stress far from the ply crack (x=5L) also agree well with thin-laminate-plate

theory.

PROCEDURE

A total of 73 [0/±45/90] , 73[0/90/±45]

s

, and 70[0g] specimens were

fabricated from the same batch of graphite/epoxy reported in Ref. 1. All

specimens were 10.2 cm long with uniform width. Quasi- isotropic laminates

were 1.27 cm wide and unidirectional specimens were 0.64 cm wide. Half of the

specimens of each type were exposed to 95 percent relative humidity at 343 K

until no additional absorbed moisture weight gain could be measured (wet or

saturated condition). The remaining specimens were dried in an oven at 338 K

until no additional moisture weight loss could be measured (dry or dehydrated

condition). The edges of four [0/±45/90]

s and four [0/90/±45]

s were polished

and replicated as outlined in Ref. 8 prior to wet and dry conditioning.

Thirty-five wet and thirty-five dry quasi- isotropic laminates (gage length

7.62 cm) were quasi-statical ly loaded in tension to fracture using a hydraulic

390
load controlled system. The cryogenic tension load fixture is shown in

Fig. 4. A special extensometer (shown in Fig. 5) was used to measure strain

over a 1.27-cm gage length with ±0.05% accuracy. Glass/epoxy tabs 1.27 cm

long were used to grip the quasi-isotropic specimens with a fine-mesh

stainless-steel interface. The stainless-steel mesh allowed thermal

contraction mismatch between graphite/epoxy specimens and glass/epoxy tabs

without causing unnecessary stress concentrations. Glass/epoxy tabs were

bonded to all [0g] specimens with epoxy adhesive. No debonding was observed

at 76 K.

All specimens with polished edges were loaded in increments of 445 N and

replicas were taken of edge damage at room temperature after each load increment.

Hence the load increment required to initiate ply cracking was recorded.

Thermal cycling after each load increment resulted in no new damage.

RESULTS AND DISCUSSION

All specimens failed at random locations between grips. There were no

apparent stress concentrations near tabs that influenced fracture strength.

Strengths reported in Table 3 compare well with results reported in Ref. 4.

Hence strengths recorded in Table 3 are minimally influenced by test

procedure. Statistical-distribution functions used to represent inherent

scatter should not be chosen a priori [9]. Here we used normal and Weibull

distribution functions and tested for "goodness of fit" using a chi-square

test [10]. All [0 Q ], [0/90/±45] and [0/±45/90] specimens fit the Weibull

distribution best with 80 percent confidence or better. Weibull distributions

imply a weak-link effect that influences fracture strength.

The largest difference between wet and dry strengths was observed for

[0/90/±45]

s laminates. Weibull strength distributions for wet and dry

[0/90/±45]

s laminates are shown in Fig. 6. Comparison of Weibull shape

391
parameters a

]A|£j>a[)pY

indl

’ cates a more dominant weak-link effect when moisture

is absorbed into the dehydrated laminate. Fracture surfaces of dry and wet

conditioned [0/90/±45]

s laminates are shown in Fig. 7 and 8. Figure 7 shows

0° ply fractures along several 90° ply cracks of a dehydrated laminate. When

absorbed moisture saturates the laminate, a single 90° ply crack dominates the

0° ply fracture shown in Fig. 8 and the Weibull mean strength decreases 8

percent.

Stress distributions near 90° ply cracks were predicted with the finiteelement model. The mechanical plate load, N , required to initiate the first

90° ply crack was observed at N = 273 kN/m. Relations (1) were used to

A calculate strains, e = 4010 ym/m and e = -964 ym/m, which constitute the

x y

mechanical load used by the finite-element model. Stresses in dehydrated

laminates loaded in tension were calculated by superposition of the mechanical

stresses with residual stresses caused by a temperature change at 76 K.

Stresses in saturated laminates loaded in tension were calculated by

superposition of additional residual stresses caused by swelling when moisture

is absorbed. Moisture weight gain was measured at 1.2 percent in the

saturated state and a temperature change of -318 K was chosen, with the

idealization of a stress-free state at 494 K (see Ref. 11). Using the above

mechanical and thermal -moisture loads together with the elastic properties

defined in Table 1, stresses near the 90° ply crack in a [0/90/±45]

s laminate

were calculated and plotted in Figs. 9 through 11.

Through-thickness variation of a in the 0° ply along x=0 is shown in

X Fig. 9. Stress in the load x direction is calculated for three cases:

MECH

1. mechanical load with no residual stresses o 2. mechanical load including

A DRY

residual thermal load o (dehydrated condition), 3. mechanical load including

A WET

residual thermal -moisture load o

x (saturated condition). At position

392
z/T = 0.753 and x=0 the largest c stresses are predicted, where

A MFfH WFT DRY

o < o < o . At position z/T = 0.766 and x = 0 the inequality of a

* X DRY WET MECH

stresses is reversed a < o < a

A A A Here we assume laminate fracture is dominated by 0° ply fracture because

the 0° plies carry the largest portion of the tensile load. Hence, the

predicted 0° ply stresses and a^

RY were compared with the mean wet and

dry fracture strengths (X^y = 425 MPa, X^

RY

= 460 MPa). In Fig. 9 stresses

averaged within the region 0.761<z/T<1.0 for saturated (WET) laminates are

larger near the crack tip than dehydrated (DRY) laminates. Hence, saturated

laminates would fracture at lower tension loads. Conversely, stresses

predicted in the region 0. 75<z/T<0. 76 indicate dehydrated laminates will

fracture at lower loads than saturated laminates. Stresses predicted in the

region 0. 75<z/T<0. 761 are located approximately one fiber diameter from the

0/90 interface. Hence the heterogeneous graphite/epoxy structure of the 0°

ply can not be modeled as a continuum within the region 0. 75<z/T<0. 76.

Stresses predicted in this region were ignored.

Delaminations were observed only at the 0/90 interface near the fracture

surface. However, delaminations at the 0/90 interface were not observed on

replicas taken prior to fracture. Since delaminations occurring at the 90°

ply crack tip will blunt the stress concentration, fiber breaks within the 0°

ply will not occur near the crack tip. In Figs. 7 and 8 we observed broken 0

ply fibers near 90° ply crack tips. Although delaminations did not influence

0° ply fracture, interlaminar stresses were plotted near the 0/90 interface i

Figs. 10 and 11. When moisture is absorbed, lower interfacial shear stress,

t , is predicted near the crack tip (see Fig. 10). Similarly, stress normal

A L~ to the 0/90 interface is lower when moisture is absorbed (see Fig. 11).

Hence, stresses leading to delaminations are decreased due to absorbed

moisture. 3 g 3
Laminates with a stacking sequence of [0/±45/90]

s were less affected by

absorbed moisture at 76 K than [0/90/±45 laminates. Absorbed moisture

increases [0/±45/90]

s laminated strength by 3.8 percent at 76 K. Fracture

surfaces of dehydrated and saturated [0/±45/90]

s laminates were similar; 0°

ply fiber fractures occurred randomly and stepwise along several 45° ply

cracks as shown in Fig. 12.

Strengths for unidirectional [0g] specimens, shown in Table 3, scattered

more than the strengths of quasi-isotropic laminates. Absorbed moisture

decreases strength 12 percent at 76 K. When unidirectional graphi te/epoxy is

bonded to 90° and 45° plies, a constraining effect on strength was observed by

Stinchcomb [12]. Strengths reported in Table 3 for [0, 90, ±45]

s specimens

are reduced to equivalent cross-sectional areas of 0° plies and compared with

[0g] strengths. This comparison is justified when the 0° ply fracture is

assumed to dominate laminate fracture strength. Unconstrained [0g] mean

strength is 23 percent lower than the fracture strength of a 0° ply

constrained by 90° and 45° plies. Detailed discussions on constraint are

given in Refs. 6 and 12.

CONCLUSIONS

Predicted dehydrated and saturated residual stresses within 0° plies near

ply cracks account for the observed differences in [0/90/±45]

s fracture

strengths, assuming 0° ply fracture dominates laminate fracture at 76 K.

Absorbed moisture swells 90° and 45° plies and changes the residual thermal 0°

ply stresses. When dehydrated laminates are saturated and loaded in tension

at 76 K, larger residual stress exists near the crack tip. Hence saturated

laminates fracture at lower tension loads. Fractographic observations confirm

this prediction, where a single 90° ply crack is observed to dominate the

394
fracture of a 0° ply. This dominance of a 90° ply crack on the 0° ply fracture

is recovered statistically, where the Wei bull shape parameter, a, indicates a

stronger weak-link effect when [0/90/±45]

s laminates are saturated. Laminate

strength is increased 8 percent when dehydrated.

Laminates with the stacking sequence [0/±45/90]

s are less affected by

absorbed moisture at 76 K than the [0/90/±45]

s laminates.

Unconstrained [0g] unidirectional strengths are lower with more scatter

when compared with constrained 0° ply strengths. Strength of unidirectional

graphite/epoxy at 76 K was improved when constrained by crossplies.

ACKNOWLEDGMENTS

This work was sponsored by NRC-NBS Postdoctoral Research Program and the

U.S. Department of Energy, Office of Fusion Energy.

395
REFERENCES

1. Kriz, R. D., "Effects of Moisture, Residual Thermal Cure Stresses, and

Mechanical Load on Damage Development in Quasi-isotropic Laminates,"

Doctoral Dissertation, Department of Engineering Science and Mechanics,

College of Engineering, Virginia Polytechnic Institute and State Universit,

Blacksburg, VA, Dec. 1979.

2. Rosen, B. W. , "Mechanics Of Composite Strengthening," in Fiber Composite

Materials, American Society for Metals, Metals Park, Ohio, 1965,

pp. 37-75.

3. Hartwig, G., "Reinforced Polymers at Low Temperatures , " in Advances in

Cryogenic Engineering- -Material

s

, Vol . 28, Edited by R. P. Reed and A. F. Clark, Plenum Press, New York, 1982, pp. 179-189.

4. Haskins, J. F. and Holmes, R. D., "Advanced Composite Design Data for

Spacecraft Structural Applications," Technical Report AFML TR-79-4208,

Air Force Materials Laboratory, Wright-Patterson Air Force Base, Ohio,

October 1979.

5. Jones, R. M. , Mechanics of Composite Materials, McGraw-Hill, New York,

1975, pp. 147-15?:

6. Talug, A., "Analysis of Stress Fields in Composite Laminates with

Interior Cracks," Doctoral Dissertation, Department of Engineering

Science and Mechanics, College of Engineering, Virginia Polytechnic

Institute and State University, Blacksburg, VA, August 1978.

7. Renieri, G. D., "Nonlinear Analysis of Laminated Fibrous Composites,"

Doctoral Dissertation, Department of Engineering Science and Mechanics,

College of Engineering, Virginia Polytechnic Institute and State

University, Blacksburg, VA, June 1976.

8. Stalnaker, D. 0. and Stinchcomb, W. W., in Composite Materials: Testing

and Design (Fifth Conference), ASTM STP 674, American Society for Testing

and Materials, Philadelphia, 1979, pp. 620-641.

9. Tracy, P. G., Rich, T. P., Bowser, R., and Tramontozzi, L. R., "On the

Statistical Nature of Fracture," International Journal of Fracture, Vol.

18, No. 4, April 1982, pp. 253-277.

10. Park, W. J., "Basic Concepts of Statistics and Their Applications in

Composite Materials," Technical Report AFML-TR- 79-4070, Air Force

Materials Laboratory, Wright-Patterson Air Force Base, Ohio, June 1979.

11. Pagano, N. J. and Hahn, H. T., in Composite Materials: Testing and

Design (Fourth Conference), ASTM STP 617, American Society for Testing

Materials, Philadelphia, 1977, pp. 317-329.

12. Stinchcomb, W. W., Reifsnider, K. L., Yeung, P., and Masters, J., in

Fatigue of Fiberous Composite Materials , ASTM STP 723, American Society

for Testing Materials, Philadelphia, 1979, pp. 320-333.

396
LIST OF TABLES

!l. Lamina properties.

12. Comparison of finite-element and laminate-plate- theory layer stresses.

3. Weibull and normal strength distributions at 76 K.

397
Table

1.

Lamina

properties

co

T“

§1CM

T- 3

CO

CM 0

CO1“0

11CM

o’

CO

UJ

ISCM

UJ

LU

0©O

CO6

00O

CO

00

<b ^ - 2

io 3

CM

CO

CO

IOgc

io

TT

00_ CM O w K

05

COo

CO^ ©a

c® o a, a a 2 © o .2 © ~

Q, IN

©

® 2 ^ 2 <0 ft <0

Js “ « © © o . V. > .2 1“ o

CO ~ QC c c n —

CS JS *

CM oa

N

CO©©O

CM X# c n

.2 E © -v, Sa |< © ” t- ® .1 *-• 3 O ©

*- oc

CO I §

3E o

CO

CM CM00

00

CM©s 0 E 1 | x i ® c ^ — .©

to o **- E” «E © ® £ o

S~

a

©eo

so©©a©

k_ao©©©©3

CO©aE©+*• Eooa

©©

CO©©

h.©>

COc

CO©©

COosc

CO

CM

T3C©co

4—o©©£«*= ©©

COo

-6c©4-» axa©

.O©0

398
Table

2.

Comparison

of

finite-element

and

laminate-plate-theory

layer

stresses

w0

oco0

jC0•*— 0a0•«-> 0cE0C0E000

•Hc

il0

CL2000w4-i CO

X

ind in o o mi cb*0)

in CO eg

y—

— • kJ0 o o o>1 o> >0a

N. CO CM CM d T— T“ in CO CM CM

•>* a> 00 CO

h** CO

CO o o m CM CM

xb

a>o 117 CO

in

co

in

in i in

m

>* 0) CM

O)

r" • • i in in 0

>*0 OO oO

a> o ©

in ini o0000C0E©03O0>OO0

O)0k_0>000000

k_

*»• 0»C0E00

399
Table

3.

Weibull

and

normal

strength

distributions

at

76K

eg CL2

C/5

*-> 08C®k_

<*«• COs

k_ >®

*oD

£0 o O o m £0 o

£0 00

m 00

© Isi- K CM q CO c© 05 in CM ID

CM CM CM CM CO CM O

00 on c©

id

CM c©

IDIs*.

T2®

u

CO CO c

.2

+-

<0

co

k. e 3 1

as leoo o CM CM CO o o

00 0)©

©

© ID

10.

K CO CO

K 00 CO

CO

O CM Is** cd

CM CO CO CM CO CM

®

QCm® \ ®

O

00

O o

CO

ID

CM CO

IDK

T" o 'O' CO

© k.®33a ! <0 o o ID ID CO K ±

©0 ID

K CM Is*. CO !£

CM

V" T— CO

© o

co

K CO cq

q cq

© Eo © so K © CM cd k_ ID CO v->k. Q

T“®

t->k. Q

®

2S

®ac® ! 2«—co > ® (0 (0 co co o

w. Q

*55' Id1 Id©

ro©

*¥* c®

r—» i—i 41 41 ***, "**>* o

00 so ID ID k. o o O O ®

»-_J f__J

©

.2,

©'*«» a

41 a41 a

CL #

400
LIST OF FIGURES

1. Graphite fiber breaks adjacent to 90° ply crack in a [0/90/±45]

graphite/epoxy laminate loaded in tension at room temperature (see

Fig. 19 in Ref. 1).

2. Quasi-isotropic laminate dimensions.

3. Finite- element grid.

4. Cryogenic tension test fixture.

5. Side-mounted extensometer.

6. Wei bull plot of wet and dry [0/90/±45]

s strengths.

7. Fracture surface of a dry [0/90/±45]

s laminate.

8. Fracture surface of a wet [0/90/±45]

s laminate.

9. Variation of o through the 0° ply thickness above the 90° ply crack tip.

Load cases: 1? Mechanical load N =273 kN/m with no residual stress

(MECH), 2. Mechanical load including residual thermal load (-318 K)

(DRY), 3. Mechanical load including residual thermal and moisture load

(+1.2 percent H

2 0) (WET).

10. Variation of t with x along surface z/T=0.766 above 0/90 interface.

Load cases are

x £he same as shown in Fig. 9.

11. Variation of a with x, along surface z/T=0.766 above 0/90 interface.

Load cases are

z the same as in Fig. 9.

12. Fracture surfaces of a wet [0/±45/90]

s laminate.

401
402

Figure

1.

Graphite

fiber

breaks

adjacent

to

90°

ply

crack

ina

[C/90/±45]

graphite/epoxy

laminate

loaded

in

tension

at

room

temperature

(see

Fig.

19

in

Ref.

1).
U, X

A

Figure 2. Quasi-isotropic laminate dimensions.

403
UNIFORM

DISPLACEMENT,

U

w, Z

Figure 3. Finite-element grid.

404

PLY

CRACK

SURFACE
Figure 4. Cryogenic tension test fixture.

405
406

Figure

5.

Side-mounted

extensometer.
Figure 6. Weibull plot of wet and dry [0/90/±45]

s strengths.

407

PROBABILITY

OF

FAILURE,

%
DRY

Figure

7.

Fracture

surface

of

a

dry

[0/90/±45]

laminate.
WET

409

Figure

8.

Fracture

surface

of

a

wet

[0/90/±45]

laminate.
100

110

120

ksi

LIZ

ire *

410

500

600

700

800

crx,

MPa

Figure

9.

Variation

ofa

x

through

the

0°

ply

thickness

above

the

90°

ply

crack

tip.

Load

cases:

1.

Mechanical

load N

=273

kN/m

with

no

residual

stress

,(M£CH),

2.

Mechanical

load

including

residual

thermal

load

(-318

K)

(DRY),

3.

Mechanical

load

inclu

ding

residual

thermal

and

moisture

load

(+1.2

percent H2

0)

(WET).
“

i gure 10. Variation of x with x along surface z/T=0.766 above 0/90 interface. Load cases

are the same aPshown in Fig. 9.

411
4 ksi

0.0 0.01 0.02

X/T

0.03 0.04

Figure 11. Variation of a with x, along surface z/T=0.766 above 0/90 interface. Load cases are the same al in Fig. 9.

412
Figure 12. Fracture surfaces of a wet [0/±45/90]

s laminate.

413

O

TENSION

ORTHOTROPIC ELASTIC CONSTANTS OF A GLASS-FIBER-CLOTH

EPOXY-MATRIX COMPOSITE

H. M. Ledbetter and M. W. Austin

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

We report the complete orthotropic-symmetry elastic-constant tensor for a laminated glass-fiber-cloth epoxy-matrix composite. Using a Marx-oscil lator

method, we determined the velocities of extensional (Young-modulus) waves in seven directions in each of the principal orthogonal planes. By a similar

method, we determined velocities of torsional waves along the three principal

orthogonal axes. From these twenty-two measurements, we deduced the ninecomponent S. . elastic-compliance matrix. By matrix inversion we computed the

C. . elastic-stiffness matrix. Unexpected results for a shear modulus and for

* J some Poisson's ratios suggest excessive fiber waviness in the fill direction.

Key Words: Composites; elastic constants; epoxy; glass fiber; Marx

oscillator; Young's modulus

415
> . .0 .....

: o-v - I
INTRODUCTION

Not long ago, Ledbetter [1] reported an experimental study of the elastic

constants of a glass-fiber-cloth epoxy-matrix laminated composite. The com- posite's construction led to orthotropic macroscopic elastic symmetry with

three principal orthogonal directions: warp, fill, and normal. Such symmetry

requires nine independent elastic constants, in contrast to two for isotropic

materials.

That study suffered two deficiencies. First, specimens were taken from

two separately prepared plates; thus, sample variability entered in an unknown

way. Second, only six elastic constants were determined from a set of thirteen sound- velocity measurements. The remaining three resulted from some assumptions concerning increased symmetry and two of the Poisson ratios.

Continued and increased interest in this material, NEMA designation

G-11CR, demanded a more thorough re-examination of the elastic constants.

Here, we report the complete nine-component elastic-constant tensor based on twenty- two independent sound-velocity measurements using a Marx-oscil lator

method.

EXPERIMENT

Both materials and measurement methods were described thoroughly previously [1]. For brevity we omit these here.

RESULTS

Figures 1 through 3 show the Young ' s-modul us variation with angle in the

three principal planes: warp-fill, fill-normal, normal-warp. Table I gives

the components of the S^. tensor together with the C^. tensor obtained by

matrix inversion. Table II recasts these as so-called engineering elastic

constants. The measured torsional moduli were 0.062, 0.056, and 0.035*

11 2

10 N/m along x^ (warp), X

2

(fill), and x^ (normal) directions, respecti vely.

417
DISCUSSION

Orthotropic-symmetry materials possess nine independent elastic con- stants. Considering the elastic compliances, the S.., these are usually taken

* J to be: S^, S^, S

33 , S 44 , Sgg, Sgg, S^> S^3

, S

23

. Young's modulus along

the i axis (i=l,2,3) is S’. j . In an arbitrary direction, Young's modulus, E,

is

:

1/E ' S

ll

~ S illl“ a

li

a lja lka

ll S

i jkl

where primes denote a co-ordinate system rotated away from x^, x

2

, x^; the

four indices on S are the full-tensor form necessary for tensor transformations; the a^. are the direction cosines between the primed and unprimed co- ordinate systems. For orthotropic symmetry. Eg. (1) becomes

Q( _ 4 Q , 4 - 4 c .oi 2 2- 2 2-

b

ll

a iril a 12

b

2 2

a 13^33 ^ a lla 12

b

12

a

1 l a 13^13

22 22 22 22 + a 12

a 13

S

23 ) + a 12

a 13

S

44 + a n

a

i 3

S

55

+ a n

a

i 2

S

66 ( 2 )

For directions in the x^-x

2 plane, Eq. (2) becomes

r i _ 4 , 4 q .2 2 / q , — \ b

11 ' a lril a 12

b

22

a lla 12 Ub 12

b 66 }

For the x 2 ~x

3

plane:

S

11

a 12

S

22

+ a 13

S 33

+ a 12

a 13^ 2S

23

+ S 44^

(3)

(4)

For the x^-Xg plane:

S 11

= a ll

S ll + a 13S 33

+ a ll

a 13 (2S

13

+ S 55 ) (5)

From Eqs. (3)-(5), by measuring E(e) in three orthogonal planes, one obtains

six elastic constants: S^, S

22 , S 33> 2S12 + Sgg, 2S

23 + S 44 , and 2S^

3 + Sgg.

To separate, for example, S^

2 from Sgg and obtain all nine S.. . requires

further measurements.

418
Torsional-modulus measurements provide the necessary information.

Torsional moduli, T., along the principal, i, axes relate to the S.. according

• J to:

T

1 = 2 ^ S 55

+ S 66^

T

2 = 2 ^ S 44

+ S 66^

T

3

= 2 ^ S 44

+ S 55^

Thus, three torsional-modulus measurements combined with E measurements versus angle in the three principal orthogonal planes lead to an unambiguous determination of all nine S.. .. By inversion of the six-by-six S.j . matrix, one obtains the C.. elastic stiffnesses:

^ J

^ C ij^6x6

=

^ S ij^6x6

Two of the three types of engineering elastic constants in Table II are readily understood because the Young modulus is simply

( 6 )

(7)

( 8 )

( 10 )

and the shear moduli are simply

G

23

C 44’ G

13

C 55’ G

12

C

66

Poisson's ratio:

(ID

S. ./S.

.

ij n

( 12 )

being the ratio of elastic compliances, transforms in a more complicated way.

Thus,

ij

= -s: ./s:

.

ij' li

where is given by Eq. (2) and

( 13 )

419
Ql _ 2 2 - 2 2 <. , 2 2 Q 3-1 i

' ^91^99^99 ' 3^o3-iq^oo ^

, 3 2

o -I 3

2

-I o • 3i

2

i Sno2 KJ j

’12 a n

a 2rn

*

^

1

3

v a

2

1

2 2. 22 22

' U 23 U 13'J 33

' vo 21u 12

' a ll

0 22 /0 12

2 2 22 7 ^ 13

' a 23

(a ll

+ a 12

)S 23

, 2/2 2 iq x 2 , 2 x 2 x c + 9i ^(doi 6o<->jS-io + 9oo(9i i *" 3 i o ) S

+ a 13

a 23a 12a 22S

44

+ a 13

a 23a lla 21

S

55

+ a 21

a 22a lla 12S

66 ( 14 )

The Young 1 s-modul us results present no surprise. We observed E^^^,

consistent with fiber packing and arrangement. The torsional moduli behave

similarly, T^T^T^, as expected.

The shear-modulus results present one surprise: G^ exceeds G^. Based

on fiber packing in the two directions, one expects the reverse. Tentatively,

we interpret this to mean that in the fill direction (x^) the fibers are more wavy than in the warp direction. That is, they possess a shorter radius of

curvature. Based on a finite-element model, Kriz [2] discusses such waviness

effects. As expected, G^ exceeds both G^ and G£^ because it involves tilting of the fiber-reinforced laminae, whereas G^ and G^ involve only their

rigid displacement, without their deformation.

The Poisson-ratio results present several surprises. First, and

differ by a factor of four! Again, one can interpret this as due mainly to

fiber waviness in the fill direction the waviness permits a larger

transverse strain in response to a tensile stress along x^. Similar reasoning

explains the large difference between and due to lower waviness, the

warp-direction fibers resist more strongly a transverse displacement, arising

from a uniaxial stress along x^ or x^. This waviness hypothesis is not con- sistent with the observed E^:E 9 ratio and the reported fiber-packing densities. Thus, we will conduct further studies before finally reporting on this

subject.

CONCLUSIONS

1. Defining x^ = warp direction, - fill direction, x^ = normal direction, we find that the Young's moduli and torsional moduli, E and T, are ordered: E^E^E^ and "^To^, as expected.

420
2.

For the shear moduli, exceeds and G

2

3

as expected.

Surprisingly , G^ exceeds G^. This can be explained by invoking excessive

waviness (short radius of curvature) for the f

i

1

1

-di rection fibers.

3.

Peculiarities in the Poisson's ratios •

of 4.0 and exceeds by a factor of 3.5 -

waviness hypothesis.

exceeds by a factor

can be understood by the same

4.

Differences in Poisson's ratios, after correction for differences in

fiber-packing density, may provide a useful tool for determining the waviness

of fibers in various directions.

ACKNOWLEDGMENT

This study was sponsored by the U.S. DoE Office of Fusion Energy.

421
REFERENCES

1. H. M. Ledbetter, to be published. Included in Materials Studies for

Magnetic Fusion Energy Applications at Low Temperatures - III, NBSIR

80-1627 (June 1980).

2. R. D. Kriz, in this report.

FIGURE CAPTIONS

[

1 ]. Directional variation of Young's modulus of G- 1 ICR composite in

11 2

warp-fill plane, in units of 10 N/m .

[2]. Di rectional

warp- normal

varia

plane

[3 J

.

Di rectional

normal -fi 1

1

varia

pi ane

tion of Young's modulus

, in units of lO^N/m^.

tion of Young's modulus

, in units of 10^N/m^.

of G- 1 ICR composite in

of G- 1 ICR composite in

422
Table I. Elastic stiffnesses, C.., and elastic compliances, S..

' J J

ij

C. . ij

(10

U

N/m

2

)

S. .

(10" Um

2

/N)

11 0.349 3.282

22 0.425 3.745

33 0.249 6.301

44 0.048 20.78

55 0.057 17.53

66 0.068 14.68

12 0.127 -0.676

13 0.090 -0.669

23 0.192 -2.641

423
Table II. Engineering elastic constants: E = Young modulus

G = shear modulus, v = Poisson ratio

11

;= 0 .,305-10

n

N/m

2 G

12

= 0 .,068.10

n

N/m

2 v 12

= 0.206

22

== 0 .,267 G

13

= 0 .,057 v 13

= 0.204

33

:= 0 .,159 G 23

= 0 .,048 V

21

= 0.180

v 23

= 0.705

V

31

= 0.106

v 32

= 0.419

424
• • • • o o o o

nid

425

Fiq.

1.

Directional

variation

of

Young's

modulus

of

G-1

ICR

composite

in

^

11

o
426
"IVWdON

427

Fig.

3.

Directional

variation

of

Young's

modulus

of

G-1

ICR

composite

in

11

2

normal-fill

plane,

in

units

of

10

N/m

.

Calculated Elastic Constants of

Composites Containing Anisotropic Fibers*

S. K. Datta

Department of Mechanical Engineering

University of Colorado

Boulder, Colorado, 80309, USA

and

H. M. Ledbetter and R. D. Kriz

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303, USA

By a wave-scattering method, we derive dispersion relationships for waves propagating perpendicul ar to continuous fibers that are oriented unidi regionally. In the long-wavelength limit one obtains relationships that predict the

composite's effective static elastic constants. We compare these relationships with others derived by energy methods to obtain upper and lower bounds

of the effective static moduli. We demonstrate them graphically by plotting

for graphi te-epoxy the predicted composite constants over the full range of

fiber volume fractions. We consider the fibers to be anisotropic, but

transversely isotropic. Under special conditions, the energy-method upper and

lower bounds compare identically with the results of this study. The static

properties are, of course, special cases of the more general dispersion

relationships. Graphs are given for nine elastic constants: axial and

transverse Young's and shear moduli, bulk and plane-strain-bulk moduli, and

three Poisson's ratios.

Key words: Bulk modulus; composites; dispersion relationship; elastic con- stants; fiber reinforcement; graphi te-epoxy ; shear modulus; wave propagation;

Young's modulus.

* Accepted for publication. International Journal of Solids and Structures

429

1. INTRODUCTION

In recent years, both theory and observation of fiber- reinforced com- posites have advanced dramatically. Reviews of this subject include those on theory by Hashin [1], Sendeckyj [2], and Walpole [3], and on experiment by

Bert [4].

In practice, two types of fiber reinforcement occur: continuous-fiber

and short-fiber (chopped-f iber) . Most studies consider only the first type,

which we consider here.

Most previous studies calculated overall (effective) static elastic

constants or bounds of such constants. Interested readers should see [5-14]

for relevant studies. Some authors considered wave propagation in fiberreinforced composites both for fibers in periodic arrays [15,16] and for

fibers distributed randomly [17-19].

Random distributions of identical, long, and parallel fibers form the

subject of the present study. We focus on the propagation of plane longitudinal and shear waves propagating perpendicular to the fibers. We use a multiple

scattering approach to obtain a dispersion relationship. This relationship

governs both propagation velocity and frequency of the incident wave. We make

three simplifying assumptions: (1) random and homogeneous fiber distribution,

(2) Lax's quasi-crystalline approximation, (3) wavelength long relative to

fiber diameter. In this long-wavelength limit, one obtains effective static

elastic constants that agree well with results obtained by other methods. For

a graphi te-epoxy composite, we consider fibers to be anisotropic, and we obtain elastic constants for all volume fractions.

2. FORMULATION

Consider a system of long circular cylindrical fibers embedded in an infinite matrix. We assume that the matrix is isotropic and homogeneous with

Lamd constants A and y and mass density p and that the fibers are transversely

isotropic with the axis of symmetry coincident with the fiber axis. We consider two problems. First, the propagation of shear waves moving perpendicular

431
to the fibers. Here the particle motion is assumed to be parallel to the

fibers. This SH-wave problem was considered in [17,19,20], The second

problem is that of propagation perpendicul ar to the fibers with the particle

motion also perpendicular to the fibers. Details of the calculation for the

first problem are given below. The second problem can be analyzed similarly,

except that the algebra is much more complicated because of its vector nature.

2.1 SHEAR WAVE POLARIZED PARALLEL TO FIBER

Let the fibers be labeled 1, 2, ..., N and let (r., e.) be the polar

t h

^ ^ coordinates of the center 0. of the i fiber. The z axis is the fiber axis.

We assume that the fibers are identical in geometry and material properties,

although for the following analysis this assumption is not necessary. In a customary notation, the fiber elastic constants are denoted by C..(i,j = 1-6).

J Since the fibers are assumed transversely isotropic there are only five

independent constants: C^, C^, C^» ancl ^44'

For the problem considered in this section, the only elastic properties

that enter the calculation are and y. Thus, the analysis follows closely

that in [16] and is a special case of that in [19], which deals with

el 1 ipti cal -cross-section inclusions.

Let an SH wave given by

u z^

=

" Yt

^’ 6 = Y//c 2’ c

?

= (!)

be incident on the medium considered. Here u represents particle displacement, c^ the shear-wave speed in the matrix, and y the circular frequency.

This incident wave will be scattered by N fibers and the scattered field

is given, using notation in [20], by:

432
( 2 ) ,(s)

N

Z i=l

( r .

)

'%i ' u, (nrJ %

r t h

where u

z is the field near the 1 scatterer and will be called the exciting

- E field [20]. The operator T ( r

)

operating on u z gives the scattered field due

"t h

* E to the i scatterer. It is well known, see [20], that Tu

z can be written

T(r.)u L (r r.) = z y [A

1 cos(n<|>.) + A sin(n<j>.)] H (bR-) (3) '%r z n n

L in v T

i

' in v T i' J n

v i' v ' n=0

where y

= 1/2, y= on 1 (nv > 0),' R. and <t>. are shown in Fig. 1, and H and J

*

1 v

i 3 ’ n n are Hankel and Bessel functions of the first kind. The coefficients A?

n and

2 A^

n depend on the boundary conditions at the surface of the fibers and are not

generally obtainable in closed form. Equation (2) can be rewritten as

. . 00 u z

= Z y n^

A

in cos ( n<f> i) + A

in

sin (^i^ H

n

( n=0

6fV

00 00 + z z u

n [a

j n cos ( n i

) + a 2

- n sin (n^.)] J^BR^ (4)

j/i n=0

where

°° ime . .

m

? 0

^[( - 1 )m(Aj,n+ m

+ (- 1 )" Aj,-n+m ) e

-ime .

.

+ (A. + (-1

)

n A. ) e

1J

v ] H (er. .) (5) j,n-m v ' j,-n-nr J nr ij' v '

433
OP L = z iU(-Dm

(A, _ - (-1)" A . _) e ime. . ij

Jn m=0

m J ,n+m J ,-n+m'

-ime .

.

+ < Aj,n- m

- (-1)" Aj.-n-m) e ^ ) ( 6 )

where

A. = i( - iA

2 ), (-1 ) PA. = ^-(A1 + A

2

) ip 2

V ip ip 7 ’ v ' i ,-p 2

V ip ip 7

t h The boundary conditions on the i scatterer are:

and

( 7 )

u -^-(u

(i) + u (s)

)| y

= C ^ I 3R^ u

z

u

z

;| R

i =a

l

44 3R

i 'R^a, ( 8 )

x. U Here u^ is the field within the i L scatterer and can be expanded in the

form:

oo = E u

n [A:J cos (n*i

) + a:

2 sin (n<j>

i )] Jp

( 6R

i

) (9)

n=0

where

0 < R. < a

.

Substituting (1), (4), and (9) into (7) and (8) yields a set of equations

for the unknowns A.

12, A. , A'. 12, and A! . The solution for A.

12

and A. can be

in’ in’ in’ in in in

written as

434
( 10 ) .1 _ ilv v .2 _ j2 v v

in nm im’ in nm im

where

m

i6x

i 1 2 a • = 2uie + Z a . , a •

i m o

j i ^ j m * l m

£ a

.

The repeated index implies summation over that index. It can be shown that

TTM = i C <5 6-1 , = iC 6 (no sum on n) (11) nm n nm lv nm n nm 2v

where

c

= Mj

n (Ba a)l?] n (ea) ' C44

J

n (ea)liJ

n (e'l )

C 44

H

n (6a)flJ

n (6

' a) ' >,J n (B' a)l?V fia)

6' = y/c^, = /C

44

/p'.

1 2 Similar expressions can be obtained for Al and Aj

n

,

Having the formal expressions for the scattered field we now proceed to

take ensemble averages. For this purpose we assume a homogeneous fiber

distribution. In this case the position of a single fiber is equally probable

within the region S, the total specimen area. Hence its distribution, p, is

uniform with density:

p<ti> \

l/s, r. e S;

o, t S.

( 12 )

V. t h Having the center of the i Lr fiber at 0^ the conditional probability density

of the j

th fiber at r. is:

435
r

l/S(l-f(r

12 )), r2

e S

(13)

where the pair correlation function f(r^) has the property:

r

(14)

approaches 0 as r^ approaches °°.

For the purpose of this study we assume that f(r.^) = 0, which is asymptotically valid for small fiber concentrations. Neglecting f(r^) im Pl ies no pair

correl ations , an assumption discussed by Bose and Mai [18]. Equation (13)

implies an isotropic distribution. Together with the quasicrystal 1 ine approximation made below in (17), this has been shown [21] to lead to optimum

bounds. By optimum we mean with respect to the Hashin-Shtri kman variational

principle [22] for isotropic distribution of spherical inclusions in an

isotropic matrix. Equation (14), often called the "wel 1 -stirred" approximation, becomes invalid at high concentrations. Nevertheless, we expect from

the discussion in [21] that the present study will lead to the general bounds

[23,24] in the 1 ong-wavel ength limit.

Taking now the ensemble average of the total field u

z we obtain:

+ r (15)

2 where n is the number density of thp fihers and C = n Tia . o

J o 436
u ( r I r.

)

Z

Vr

v.

1 ' - u o

e

l 6x

"0/

dr

.

T(r .

)

u^(r I r

.

r .

)

(16)

| r -r | > 2a.

%i

It is seen that the first partial average of the exciting field at r . depends

on the second partial average of the same. An equation for the second partial

average will depend on the third, and so on.

This hierarchy is terminated here by making the quasi -crystal 1 ine

approximation that requires

<u ( r I r .

,

z

v ^ 1 %j

’ r .

)

= <u ( r I r .

)

>. (17)

This makes (16) an integral equation for <u^( r | jr ^ ) >

To solve this equation it may be assumed that

<u^(r|r.) > = z y [a?" cos(n<j>.) + b^ sin(n<j>.)] J (b R • ) . (18)

n= Q

M nL in v ' in v y i' J n VP v '

Then from (3)

co

T(jCu) < u f(JCki> > = Z Q

y

n [Ajn cos ( n ^j) + A

j n sin ( n <f>j)] H

n (eRj)- (19)

Thus we obtain:

IP

d

l6x

i = 2u i p

e

o n^ fdr. z iy { (T

°J -0 n=0

2 n

IV

nq

T

2v

) nq

> 2a

437
-i(p-n)e. . i(p+n)e..

x [H (er. .)e J + ( -1

)

n

H e

1J

L ] p-n

x v ' p+n J

o

i (p-n)e. . i (p+n)e.

.

+ (T

lv

+ i t

2v

) [He 1J + (-l)

n H , e

1J ]} a v

v nq nq' L p-n

'

' p+n J jq

( 20 )

ip

_ 00 n /dr . z ip { (

T

V ^ n=0

2 n lvnq

.o -i (p-n)e . .

i T

2v

) [H e

1J

nq' L p-n

r .-r . > 2a

\J -vi 1

i (p+n)e. i (p+n)e .

.

(-1)" H x

e

r ' 1j

] - (T

lv

+ i T

2u

) [H e " "

1J

p+n nq nq' u p-n

n -i ( p+n ) e. . (-1)" H p+n

e ,J ])

.J

• ( 21 )

Equations (20) and (21) define a pair of integral equations in the unknowns

1 2

a.. and a^ . If it is assumed that an effective plane wave is propagating

through the medium, then a solution for a^ may be assumed in the form

ie*x. a-n = X n e ip p

( 22 )

where y/B* is the wave velocity of the plane wave.

The coefficients Xp form a set of linear homogeneous algebraic equations

These are:

438
. 2irn a

00 Y 1 __ 0 v

• P"H A 4-A \ Xp "

" b*

2 -6 2 n=0

yn

Tpq Xq Ap “ n Ap+n

0 2nn a

00 „ AO

X

2 = E y i P_n

T

2v

X V(A + A ) p 3*

2 -B2 n=0

n nq q p

“ n p+n

where

4m tfl H

m'

26a )] J

m< 26*a > ' & J m

< 2B* a >] Hm< 2Ba >

( 23 )

(24)

(25)

1 2 Because of relationships (11), X^ and Xp are uncoupled. However, for noncircular-cross-section fibers they are coupled. The dispersion equation is

1 2 obtained by equating the coefficients of Xp and Xp to zero. For an arbitrary

wavenumber the dispersion equation is rather complicated. We can approximate

this equation by assuming that 6a is small; this means the wavelength is long

compared to the fiber diameter. For small 6a and 6

1 a the following relationships can be obtained, which are correct to order (6a) :

C

Q

- (6a)

2 (p/p

1 - 1), C

y ~ 2 ^44 - y C

44

+ u (26)

C = 0 for m > 2 m —

and

a - /2i, / 6* \ A m <na>

} m

(27)

Substituting (26) into (11), and then into (23), and using (27) one obtains:

439
2 7T n a o

oo (28) P y

1 - 2 r2 n=0

,2 \ .-n

u n

( ^> 1 X

1 C

n n {( ?lL\ |P" n <£>

p+n

}.

Because of (26) the infinite series can be terminated after n = 1. This

yields two simultaneous homogeneous equations in Xq and xj. Equating the

determinant of the coefficients to zero gives the dispersion relationship:

[1 + C ( p 7p

- 1)]

1 - C(m - 1 )/(m + 1

)

1 + C(m - 1 )/ (m + 1)

(29)

where m = C^/y and C is the fiber volume fraction.

Defining the average density as

(30)

(31)

where y^y is the effective axial shear modulus, one obtains:

p* - p[l + C(p'/p - 1)]

and

6* = y/c

2 *, c

2 * - /u LT

/p*

_LT 1 + C(m - 1 )/ (m + 1)

y m

1 - C(m - 1 )/ (m + 1)

( 32 )

where m is the ratio of fiber and matrix shear moduli ( p

j/

y

m

)

440
Equation (32) was obtained by Hashin and Rosen [25] using the composite

cylinder assemblage (CCA) model. The expression for obtained from (32)

coincides with the lower (upper) bound when the fiber shear modulus, y^y, is

larger (smaller) than the matrix shear modulus, y .

2.2 LONGITUDINAL AND SHEAR WAVES POLARIZED PERPENDICULAR TO FIBER

Propagation of longitudinal and shear waves polarized transverse to the

fibers was studied in [18] using the same approach outlined above. It was

shown that if 6a and aa were assumed small then the dispersion relationships

reduce to effective longitudinal and shear moduli transverse to the fibers.

These are:

C n 1

- CP

2 (1 - 6 2

/a

2

) - 2C

2

P qP

2

A m

+ 2ym (1 + CP

Q

) [1 + CP

2 (1 + 6 2 /a

2 )]

(33)

TT

= 1 +

2 C(y

TT

‘ y m

} (Am

2 y (A + 2 y ) + (1 - C) (

A

m

v m m m IV+ 3M

m>

(p TT

- u (34)

nr

where

p

. K

i

- (xm

+ um ) _

0 K

f

+ ' 2 ‘7 tt

tx

m

TT

• 3 y ) m

' '

, m

)

pm

)

a

2

= Y

2

/c

2

, Cj = /(Xm+ 2 p m

)/p.

The effective transverse plane-strain bulk modulus, Ky, can be obtained from

equations (33) and (34) as:

441
( 35 )

c(k; - a - p ) KT

= A + y + (A + 2y ) 7 r \ v ,—

7^-7 —"7'

i—1' r\ T m m m m (1v - C)IC + CA + (1 + C) y T m

v '

where K-j. is the transverse bulk fiber modulus and the matrix Lamd constants

are A and y . m m

Equation (35) is the same as derived by Hashin and Rosen [25] using the

CCA model. However, the CCA model does not yield an expression for the

effective transverse shear modulus. Instead, bounds for the transverse shear

modulus are noted. Expression (34) for y^ is identical to the general lower

(upper) bound derived for arbitrary phase geometry [23,24] if y-p. and K-j- are

larger (lower) than y and K , respectively, where K = A + y is the planem m

r m m m strain bulk modulus of the matrix.

Knowing the expression for one can predict the effective longitudinal

Young's modulus and Poisson's ratio using relationships derived by Hill [23].

When both fiber and matrix phases are transversely isotropic, Hill shows that:

4C( 1 - C) (v* - v r

E

L " E m

(1 ' c ) + E

L C +

(1 - C)/K| + C/K + 1/y

T m

M m

(36)

C ( 1 - C) (v' - vm

) (1/K - 1 / K'

)

V

LT

= v m

(1 " C ) + V LT

C +

(1v - C)/K{ ' + C/K + 1/y

T m

M m

(37)

where E^ and are the axial Young's modulus and axial Poisson's ratio of

the fiber. The quantities subscripted with m designate the isotropic matrix

properties

.

442
For a unidirectional fibrous composite with transversely isotropic

fibers, this completes the derivation of the five independent effective

elastic properties: Ky, E^» v^y, Vyy» u

L y. Other elastic properties of

physical significance can be calculated from the five effective properties.

For instance, assuming isotropy in the transverse plane, T, the Poisson's

ratio, Vyy, and the Young's modulus, Ey, can be calculated from the isotropic

relationships. Thus,

4K

E I

- t

m tt

TT

(38)

K

T

" ^TT

V

TT

=

Kt

+ *i, TT

(39)

where

*

= 1 4 Vlt

e

l

The remaining Poisson's ratio, Vy

L

, is calculated from the relationship

obtained from the elastic-compliance symmetry condition:

U

TL

V

LT VE L- (40)

If the composite behaves transversely isotropic with the unique symmetry

axis along x^, then the elastic-stiffness constants, C^. . , can be computed from

C 11C 22KT

+ y TT

(41)

443
( 42 ) C 33

E

L + 2v

LT

C 13

C 44

= C 55

= y LT ( 43 )

C 12

= Kt

- y TT (44)

C 13

= C 23= 2v

LT

KT

( 45 ^

C 66

=

^ C ll

" C 12^ /2 = y TT. ^ 46 ^

One can also calculate the bulk modulus, B, from

B-7. * - /iijj * {E

L + + \T Kt + 2 >^9 ' < 47 >

3. RESULTS AND DISCUSSION

The effect of fiber volume fraction, C, on the predicted elastic constants can now be demonstrated given the isotropic matrix properties and the

transversely isotropic fiber properties. The matrix and fiber elastic constants, listed in Table 1, were obtained from [26] where the fiber elastic

properties were extrapolated from experiments of Dean and Turner [27] using

the transversely isotropic relations derived by Hashin [1], This extrapolation of fiber properties was done with the correction 6 2 = K-j-/ ( K-j. + 2y-j- T

)

noted in [28]. The elastic properties (v-j-y, v-y. , v^, y-j-y, Ey, U|_y> E|_» ^y>

B) are shown in Figs. 2-5.

As noted above, the curves for K-y, y^, and y^ coincide with the general

lower bounds [28] for the particular constituent properties considered here.

444
Of the nine elastic constants shown in Figs. 2-5, best fits a simple rule

of mixtures. (The correction term in (36) can be neglected.) Both B and

come close to fitting this rule; Uyy, Ky, and Ey do not differ dramatically

from linearity; but Uyy, Vyy, and Vy^ do so differ. Except for v-p and Ey all

deviations from linearity are negative, that is, the curves are concave. The

Vyy curve is convex. The Ey curve is convex at low fiber fractions and

concave at high. This reflects the abrupt rise of Vyy at low fiber fractions

combined with the slow, steady rise of pyy. For those constants that differ

strongly from a rule of mixtures, both Vyy and Vy^ are f iber-property dominated

while y^y Tends to be matrix-property dominated. The latter represents, of

course, shear in the fiber direction.

4. CONCLUSIONS

Using a multiple-scattering approach, effective elastic constants of a

graphite-fiber-reinforced epoxy composite were derived in this study. Graphite

fibers were assumed to be anisotropic, but transversely isotropic. It was

shown that the composite can be characterized as a transversely isotropic

medium. All five elastic constants characterizing the composite were calculated. These elastic constants coincide with the general lower bounds of

those obtained in [1,23,24] using energy methods.

ACKNOWLEDGMENTS

Composites studies of H.M.L. and R.D.K. are supported mainly by the

U.S. Department of Energy, Office of Fusion Energy.

445
REFERENCES

1. Z. Hashin, Theory of Fiber- reinforced materials, NASA Report CR- 1974

(1972).

2. G. P. Sendeckyj, Elastic behavior of composites, in Composite Materials ,

Volume 2 , Academic (1974), pp. 45-83.

3. L. J. Walpole, Elastic behavior of composite materials: theoretical

fundamentals, in Advances in Applied Mechanics , Volume 21 , Academic

(1981), pp. 169-242.

4. C. W. Bert, Experimental characterization of composites, in Compos i te

Material

s

, Volume 8 , Academic (1975), pp. 73-133.

5. J. J. Hermans, The elastic properties of fiber reinforced materials when

the fibers are aligned, Proc. Koninklijke Netherlandse Akademie van

Wetenschappen 70B , 1-9 (1967).

6. J. M. Whitney, Elastic Moduli of unidirectional composites with anisotropic filaments, J. Compos. Mater. _1> 188-193 (1967).

7. J. M. Bloom and D. F. Adams, Axial loading of a unidirectional composite -- anisotropic filaments, J. Compos. Mater. _3, 186-188 (1969).

8. M. D. Heaton, A calculation of the elastic constants of a unidirectional

composite containing transversely isotropic fibres, J. Phys. D, _3>

672-676 (1970).

9. C. H. Chen and S. Cheng, Mechanical properties of anisotropic fiber

reinforced composites, J. App 1 . Mech. 37, 186-189 ( 1970).

10. C. H. Chen, Rectangular or square array fibrous composite with anisotropic or isotropic filament, J. Appl . Mech. 38, 710-712 (1971).

11. E. Behrens, Elastic constants of f iber-reinforced composites with transversely isotropic constituents, J. Appl. Mech. 38, 1062-1065 (1971).

446
12. F. Ya. Bulavs, Ya. V. Auzukalns, and A. M. Skudra, Deformation characteristics of plastics reinforced with high-modulus anisotropic fibers,

Polymer Mech. £3, 545-552 ( 1972).

13. H. T. Yeh, Variational bounds of unidirectional fiber reinforced composites, J. Appl . Phys. 44, 662-665 (1973).

14. G. A. Ermakov, A. G. Fokin, and T. D. Shermergor, Effective moduli of

elasticity for materials reinforced by anisotropic fibers, Mech. Solids

9, 98-104 (1974).

15. J. D. Achenbach, Generalized continuum theories for directionally reinforced solids, Arch. Mech. 28, 257-278 (1976).

16. M. Hlavacek, A continuum theory of fiber- reinforced composites, Int. J.

Solids Structures 9_, 1075-1085 (1973).

17. S. K. Bose and A. K. Mai, Longitudinal shear waves in a fiber- reinforced

composite, Int. J. Solids Structures _9, 1075- 1085 ( 1973).

18. S. K. Bose and A. K. Mai, Elastic waves in a f iber-reinforced composite,

J. Mech. Phys. Solids 22, 217-229 (1974).

19. S. K. Datta, Propagation of SH-waves through a fiber-reinforced composite -- elliptic cylindrical fibers, J. Appl. Mech. 42_, 165- 175 (1975).

20. S. K. Datta, Scattering by a random distribution of inclusions and

effective elastic properties, in Continuum Model of Discrete Systems ,

Solid Mechanics Study Series No. 12 , University of Waterloo Press (1978),

pp. 111-127.

21. J. R. Willis, Variational and related methods for the overall properties

of composites, in Advances in Applied Mechanics , Volume 21 , Academic

(1981), pp. 1-78.

22. Z. Hashin and S. Shtrikman, On some variational principles in anisotropic

and nonhomogeneous elasticity, J. Mech. Phys. Solids H), 335-342 (1964).

447
23. R. Hill, Theory of mechanical properties of f ibre-strengthened materials - I. elastic behaviour, J. Mech. Phys. solids _1_2, 199-212 ( 1964).

24. Z. Hashin, On elastic behaviour of fibre- reinforced materials of arbitrary phase geometry, J. Mech. Phys. Solids J_3, 119- 134 ( 1965).

25. Z. Hashin and B. W. Rosen, The elastic moduli of fiber- reinforced materials, J. Appl . Mech. 31, 223-232 (1964).

26. R. D. Kriz and W. W. Stinchcomb, Elastic moduli of transversely isotropic

graphite fibers and their composites, J. Exper. Mech. 19_, 41-49 (1979).

27. G. D. Dean and P. Turner, The elastic properties of carbon fibers and

their composites, Composites 4_, 174- 180 ( 1973).

28. Z. Hashin, Analysis of properties of fiber composites with anisotropic

constituents, J. Appl. Mech. 46, 546-550 (1979).

448
List of Figures

Fig. 1.

Fig. 2.

Fig. 3.

Fig. 4.

Fig. 5.

The coordinate system.

Predicted axial Young's and shear moduli.

Predicted transverse Young's and shear moduli.

Predicted bulk modulus and plane-strain bulk modulus.

Predicted Poisson's ratios.

449
450

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1.

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system.
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TECHNOLOGY TRANSFER

455

TECHNOLOGY TRANSFER PROGRAM

LEADER: N. J. Simon, NBS

STAFF: P. 0. Chavez, NBS

OBJECTIVES:

(1) To hold a biennial workshop in the autumn of alternate years.

Major topics include design overviews of fusion energy projects

with respect to low- temperature material needs and reviews and data

on structural alloys, composites, and weldments.

(2) To produce and distribute an annual report as well as monthly highlight reports to OFE/DoE.

(3) To prepare handbook pages and supporting documentation on lowtemperature mechanical and physical properties of structural materials for superconducting magnets. The pages are to be distributed

to workshop attendees and others in the low- temperature materials

applications community. The pages will also be presented for inclusion in the "Materials Handbook for Fusion Energy Systems."

RESEARCH PAST YEAR (1982):

(1) "Materials Studies for Magnetic Fusion Energy Appl ications- V,"

NBSIR 82-1667, 490 pages (May 1982), editors R. P. Reed and N. J.

Simon, has been sent to the Technology Information Center, ORNL,

and approximately one hundred copies have been requested by other

interested persons.

(2) Handbook pages and supporting documentation covering mechanical and

physical properties of AISI 316 stainless alloys and Fe-21Cr-6Ni9Mn have been distributed in preliminary draft form to interested

persons. These pages (410 in total) have also been submitted to

the "Materials Handbook for Fusion Energy Systems."

(3) Efforts were undertaken to encourage international cooperation in

research to develop radiation-resistant laminates to meet requirements for the Engineering Test Reactor (M. B. Kasen, NBS). See

discussion and letter on following pages.

RESEARCH THIS YEAR (1983):

(1) Current plans are to hold the biennial 1 ow- temperature materials

workshop again in October 1983.

(2) "Materials Studies for Magnetic Fusion Energy Appl ications-VI ,"

NBSIR 83-1690, 466 pages (1983), editors R. P. Reed and N. J.

Simon, is the current edition of the annual report.

(3) The AISI 316 handbook pages are currently in the review process for

"Materials Handbook for Fusion Energy Systems," and the Fe-21Cr6Ni-9Mn handbook pages will enter the review process later this

year.

(4) Preparation of AISI 304N and 304LN handbook pages will be completed

this year. Preparation of handbook pages on the aluminum alloys

2219 and 5083 will be started.

(5) Efforts to develop international cooperation in research to develop

radiation-resistant laminates will continue. See discussion following (M. B. Kasen, NBS).

457
INTERNATIONAL COOPERATION IN RESEARCH TO DEVELOP

RADIATION-RESISTANT LAMINATES

During the past year, an independent program entitled "Development of

Organic Insulators for Fusion Magnets" has been established under the DoE

Office of Magnetic Fusion Energy. The program participants (LANL, ORNL, MIT,

and NBS) have the task of developing radiation-resistant laminates suitable

for use in the Engineering Test Reactor (ETR), for which construction is

expected to begin in 1988. A complimentary program to refine cryogenic

nonmetal lie test specimens and techniques has also been independently es- tablished under a subcommittee of the Metals Properties Council. A common element in both programs is the desire to develop international cooperation

in the research efforts. This was informally begun by conversations between

interested parties from the United States, Japan, Germany, and England during

the 1982 ICEC9/ICMC conference in Kobe, Japan. The attached letter outlining

the United States programs and proposing an initial step in implementing such

cooperation is an outgrowth of that meeting.

458
December 2, 1982

this letter sent to names checked on the attached list

Dear

There is an expanding awareness of the need to develop nonmetal lie

insulators suitable for use in superconducting magnets for magnetic fusion

energy applications. The required research is complicated by the need to

subject specimens to high neutron fluences at 4 K, and by the difficulty in

evaluating insulator degradation in such an environment. Few facilities are available for conducting such work, and experimental costs are high. It is

therefore recognized that close cooperation between laboratories and individuals

pursuing such research would be mutually beneficial.

This subject was informally discussed during the 1982 ICEC9/ICMC conference

in Kobe, Japan, in a group composed of Prof. Hitoshi Yamaoka (Kyoto U.), Dr.

Shigihiro Nishijima (Osaka U.), Dr. David Evans (Rutherford and Appleton

Labs., UK), Dr. Gunther Hartwig ( Kernforschungszentrum, W. Ger.), and myself.

It was agreed that the first effort should strive for agreement on specimen

design to facilitate intercomparison of results of irradiation programs. I agreed to coordinate the U.S. participation in this effort. However, I have

delayed further communication until the present time because we have been in

the process of structuring several programs having a bearing on the subject.

This has now been done.

The purpose of this letter is to summarize the objectives of these programs

for you and others who have an interest in such research, and to solicit

suggestions for ways in which mutually beneficial cooperative efforts might be

implemented between these programs and your activities.

We have now begun a six year program to develop organic insulators

capable of meeting the anticipated requirements of the magnetic fusion energy

Engineering Test Reactor (ETR) which is anticipated to enter the design and

construction phase about 1988. This program is sponsored by the Office of

Fusion Energy of the U. S. Department of Energy (DoE). It is the outcome of a series of meetings and workshops held over the last few years, the most recent

being on September 8-9 of this year. Direct participants in this program are:

Dr. Marvin Cohen, DoE, Contract Monitor

Dr. Ralph R. Coltman, Oak Ridge National Laboratory (ORNL)

Dr. George F. Hurley, Los Alamos National Laboratory (LANL)

Dr. Herbert Becker, Plasma Fusion Center, Massachusetts Institute of

Technology (MIT)

Dr. Maurice B. Kasen, National Bureau of Standards (NBS)

459
As presently structured, the program will initially address several

questions that have arisen concerning insulator behavior under simultaneous

stress and irradiation. The next objective would be to identify and standardize the best insulator materials based on the results of irradiation test

programs. Eventually, we believe that the suitability of such materials must

be verified by testing magnet components in a simulated magnet environment.

Some questions to be initially addressed are:

1. Does the majority of the structural damage to insulators occur during

irradiation at 4 K or during warmup following irradiation?

2. How does degradation under neutron irradiation at 4 K differ from that

produced by gamma irradiation?

3. Is it necessary to evaluate insulator performance under combined

stress and irradiation at 4 K?

4. What is the effect of different types of structural damage on electrical insulating performance at 4 K?

5. Will the allowable radiation fluence be limited by damage to the

insulators or to the superconducting stabilizer?

The ORNL studies will continue their evaluation of gamma irradiation

effects using their low temperature irradiation facility (LTIF) to assess the

effect of warmup on damage accumulation in epoxy. ORNL will also continue

studies of the effect of irradiation on the magnetoresistance of copper

stabilizing materials. The LANL studies are addressing the question of the

relative effect of neutron and gamma irradiation on insulators. Materials

similar to those studied by ORNL are being irradiated at the IPNS-1 accelerator base neutron source at the Argonne National Laboratory (ANL). The MIT

studies are concentrating on the interaction between the geometrical aspects

of insulator design and the properties of the insulators themselves. The

current interest is in characterizing the performance of insulators under

compression normal to the reinforcement direction. My work at NBS will

involve selection and characterization of materials required by the other

participants for their test programs, cooperating with the others in postirradiation failure analysis, and in stimulating international cooperation in

the overall effort. I am pleased to report that the value of such cooperation

has been recognized as an integral part of the program.

A second and somewhat complimentary program addresses the need to improve

and standardize cryogenic test methods. This effort is being organized by a subcommittee of the Metals Properties Council (both nonmetallic and metallic

materials are being addressed). Two organizational meetings have thus far

been held, resulting in the establishment of a nonmetallics committee consisting of Dr. Herbert Becker (MIT), Dr. James Gauchel (Owens-Corning Research

Laboratory), and myself. Dr. Gauchel has elected to work on development of a flexural test specimen optimized for use in material performance screening.

Dr. Becker has chosen to address the development of an improved technique for

measuring interlaminar shear strength.

460
We would welcome your active participation in any aspect of either of

these programs. However, as we discussed at Kobe, perhaps the best way to

begin a cooperative effort would be to investigate the possibility of establishing commonality in specimens used to screen the effects of cryogenic

irradiation so as to facilitate meaningful exchange of test data between

laboratories. As suggested in Kobe, an exchange of test specimens might be a proper first step. On our part. Dr. Gauchel has offered to provide samples of

his proposed specimen configuration that can be forwarded to you. His

specimen looks very interesting in that it can be easily and inexpensively

manufactured in any laboratory, any desired fiber content can be easily and

precisely obtained, its size lends itself to efficient screening under

cryogenic irradiation, and I understand that the configuration facilitates

analysis of test results in terms of structural factors. We would, of course,

solicit your suggestions for alternate configurations.

I also suggest that it might be beneficial for each of us to have a more

detailed understanding of the ongoing and planned activities of the various

individuals working on this problem. If each of you would take the time to

prepare a brief summary of your activities and plans, I would be pleased to

combine them into a single document and distribute it to all parties.

The attached list includes those individuals who to my knowledge are working in this area or have indicated an interest in cooperative efforts. I shall be sending a copy of this letter to each of them. Please let me know if

there are others who you believe should be added to the list.

Sincerely,

M. B. Kasen

Fracture and Deformation Division

461
Distribution List

Dr. Saburo Takamura

Solid State Physics Laboratory No. 2

Japan Atomic Energy Research Laboratory

Tokai Research Establishment

Tokai-Mura, Ibaraki-ken 319-11

Japan

Dr. Kenzo Kadotani

Senior Researcher, Nonmetallics

Hitachi Research Laboratory

Hitachi Ltd.

4026 Kuji-chuo, Hitachi-shi

Ibaraki-ken 319-12

Japan

Prof. Hitoshi Yamaoka

Research Reactor Institute

Kyoto University

Kumatori, Osaka-shi 590-04

Japan

Dr. Toichi Okada

Institute of Science and Industrial Research

Osaka University

Ibaraki, Osaka 567

Japan

Dr. David Evans

Chemical Technology Division R34

Rutherford Appleton Laboratory

Chilton, Didcot

Oxon 0X11 OQX

England

Dr. Gunther Hartwig

Institut fur Technische Physik

Kernforschungszentrum, Karlsruhe

Postfache 3640

7500 Karlsruhe

West Germany

Dr. George F. Hurley M.S. 546

Los Alamos National Laboratory

P. 0. Box 1663

Los Alamos, NM 87545

Dr. Ralph R. Coltman

Building 3115

P. 0. Box X Oak Ridge National Laboratory

Oak Ridge, TN 37840

462
Dr. Herbert Becker

MIT Plasma Fusion Center

167 Albany Street

Cambridge, MA 02139

Dr. Marvin Cohen

Office of Fusion Energy

Materials and Radiation Branch M.S. 6-234

U. S. Department of Energy

Washington, DC 20545

Dr. James Gauche!

Owens-Corning Fiberglas Research Center

P. 0. Box 415

Granville, HH 43023

Dr. Ursula Kreibich

Ciba-Geigy Research Laboratory

Ciba-Geigy AG

CH-4002 Basel

Switzerland

Dr. Maurice B. Kasen

Fracture and Deformation Division 562

National Bureau of Standards

Boulder, CO 80303

463
ACKNOWLEDGMENT

The editors wish to express their appreciation to Diane Bieber, JoAnne

Wilken, and Peggy Chavez for their work in preparing the manuscript.

Marilyn Stieg removed many of the inconsistencies, grammatical errors,

and awkward usages of language. The editors are responsible for those that

rema i n

.

464
NBS-114A (REV. 2-BC) U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO.

NBSIR 83-1690

2. Performing Organ. Report No. 3. Publication Date May 1983

4. TITLE AND SUBTITLE

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures—VI

5. AUTHOR(S)

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

6. PERFORMING ORGANIZATION (If joint or other than NBS. see instructions) 7. Contract/Grant No. NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City. State. ZIP)

Department of Energy

Office of Fusion Energy

Washington, DC 20545

10. SUPPLEMENTARY NOTES

Related earlier reports: NBSIR 81-1645, NBSIR 80-1627, NBSIR 79-1609, NBSIR 78-884,

NBSIR 82-1667

2 Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bi bl iography or literature survey, mention it here)

This report contains results of a research program to produce material property data

that will facilitate design and development of cryogenic structures for the superconducting magnets of magnetic fusion energy power plants and prototypes. The program was conceived and developed jointly by the staffs of the National Bureau of Standards and

the Office of Fusion Energy of the Department of Energy; it is managed by NBS and

sponsored by DoE. Research is conducted at NBS and at various other laboratories

through subcontracts with NBS.

The reports presented here summarize the sixth year of work on the low-temperature

materials research program. Highlights of the results are presented first. Research

results are given for the four main program areas: structural alloys, weldments and

castings, nonmetallics , and technology transfer. Objectives, approaches, and achievements are summarized in an introduction to each program area. The major portion of the program has been the evaluation of the low-temperature mechanical and physical properties of stainless steel base metals, welds, and castings, with

particular emphasis on the nitrogen-strenthened stainless steels. Developmental

steels with manganese additions are also under investigation. Codes and standards

development received increased emphasis this year, especially as related to lowtemperature property measurements. Work has been done of the standardization of industrial laminates for low-temperature applications and on the measurement of nonmetallic

composite properties at cryogenic temperatures. Modeling of the elastic and fracture

behavior of woven fiber-reinforced composites has been initiated this year.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolon s) alloys; austenitic steels; castings; composites; cryogenic properties; low temperatures;

mechanical properties; nonmetallics; elastic properties; stainless steels; structural

alloys; weldments.

13. AVAILABILITY 14. NO. OF

PRINTED PAGES

|X | Unlimited

| | For Official Distribution. Do Not Release to NTIS 436

2 Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C.

20402. 15. Price

Order From National Technical Information Service (NTIS), Springfield, VA. 22161 $32.50

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