NBSIR 79-1609

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A111Q4 MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES

X

TEMPERATURE

STRESS INTENSITY RANGE

TO

DEPARTMENT OF ENERGY

OFFICE OF FUSION ENERGY

WASHINGTON, D.C. 20545

•QC

100

.056

79-1609

BY

FRACTURE AND DEFORMATION DIVISION

AND

THERMOPHYSICAL PROPERTIES DIVISION

NATIONAL BUREAU OF STANDARDS

BOULDER, CO 80303

NBSIR 79-1609 ;EP 1 T \m UC 20 b,c

TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES

Edited By

F.R. Fickett

Thermophysical Properties Division

and

R.P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, CD 80303

June 1979

Sponsored By

Department of Energy

Dffice of Fusion Energy

Washington, D.C. 20545

U.S. DEPARTMENT OF COMMERCE, Juanita M. Kreps, Secretary

Sidney Harman, Under Secretary

Jordan J. Baruch, Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS. Ernest Ambler, Director

CONTENTS

Page

SUMMARY V

ORGANIZATIONAL CONTACTS vii

INTRODUCTION 1

PROGRAM DESCRIPTION 3

HIGHLIGHTS OF RESULTS 7

ALLOYS FOR LOW TEMPERATURE SERVICE 15

Testing of 2219-T87 Aluminum Alloy at 4 K

- Martin Marietta 17

Low Cycle Fatigue Testing of 304LN Stainless

Steel and 5083-0 Aluminum Alloys - Martin Marietta ... 33

Fracture and Strength Properties of Selected

Austenitic Stainless Steels at Cryogenic

Temperatures - NBS 79

Heating and Strain Rate Effects in AISI 304L

Stainless Steel at 4 K - NBS 123

Low Temperature Deformation of Fe-18Cr-8Ni

Steel - NBS 149

PHYSICAL PROPERTIES OF STAINLESS STEELS 173

Elastic-Constant Variability in Stainless

Steel 304 - NBS 175

Low Temperature Specific Heat of Two Stainless

Steels - Wichita State Univ./NBS 197

Magnetic Susceptibility and Magnetization Studies

of Some Commercial Austenitic Stainless Steels - Battel le Columbus 203

Some Magnetic Properties of the "Nonmagnetic"

Stainless Steels - NBS 281

STAINLESS STEEL WELDING 295

Evaluation of Stainless Steel Weld Metals at

Cryogenic Temperatures - NBS 297

Metallurgical Factors Affecting the Toughness of 31 6L

SMA Weldments at Cryogenic Temperatures - NBS 313

Microstructure of Specific Stainless Steel

Weldments being Evaluated for Cryogenic

Service - Colorado School of Mines 353

NONMETALLICS FOR MAGNET STRUCTURES 401

Nonmetallics for Magnet Systems - NBS 403

TECHNOLOGY TRANSFER 453

Structural Alloys for Superconducting Magnets

in Fusion Energy Systems - NBS 455

Handbook Preparation - MPDC/NBS 507

Vail Workshop 515

i i i

SUMMARY

The reports presented here summarize the work of the low temperature

materials research project for the second year of the program. The various

projects are outlined and the research results are presented. The major

thrust of the measurements has been the evaluation of the low temperature

mechanical and physical properties of stainless steel base metal and welds,

with particular emphasis on the nitrogen strengthened stainless steels.

Aluminum alloys have received some consideration also. Work has also been

done on the production of nonmetal 1 ics , primarily industrial laminates, for

low temperature applications and on the measurement of their properties at

cryogenic temperatures. The second NBS/DOE Vail workshop was held in October 1978. A brief description is given of that program.

Note : The use of trade names of commercial materials is essential to the

proper understanding of the work presented. This use in no way implies

approval, endorsement or recommendation by the National Bureau of Standards.

Contributions to this document by non-NBS authors have not been subjected to

the usual NBS editorial review process.

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ORGAMIZA-^TONAL CONTACTS

Listed here are the people responsible for the various major aspects

of the program so that those with specific technical questions may contact

the individual directly.

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

Program Monitor E. N. C. Dalder (301) 353-4964

(FTS) 233-4964

National Bureau of Standards, Boulder, Colorado 80303

Program Manager

Program Coordinator

and Magnetic Properties

Weldi ng

Elastic Properties

Nonmetal lies

R. P. Reed

F. R. Fickett

H. I. McHenry

H. M. Ledbetter

M. B. Kasen

(303)499-1000 X3870

(FTS)323-1000 X3785

X3268

X3443

X3558

Martin Marietta Aerospace, Denver Division, P.O. Box 31, Denver, Colorado 80201

Fatigue, Fracture F. R. Schwartzberg (303) 973-3225

J. A. Shepic

Mechanical Properties Data Center, 13919 West Bay Shore Drive, Traverse City,

Michigan 49684

Program Manager R. C. Braden (616) 947-4500

Data Evaluation T. Moore

Battelle Columbus Laboratories, 505 King Avenue, Columbus, Ohio 43201

Project Manager E. W. Col lings (614) 424-4355

(FTS) 976-4355

INTRODUCTION

This report contains results of a research program to produce

materials properties data to facilitate design and development of cryogenic

structures for superconducting magnets, required for development of

magnetic fusion energy. 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. This program, sponsored by DOE, is

managed by NBS. Research is conducted at NBS and at various other

laboratories , through subcontracts with NBS.
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PROGRAM DESCRIPTION

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

construction and safe operation of low temperature magnetic fusion

energy (MFE) systems, especially superconducting magnets, through effective

materials research and materials technology transfer. The specific

steps taken to achieve this objective are: (1) evaluation of low temperature materials research needs specific to MFE devices; (2) development

and monitoring of a research program to acquire the necessary data; and

(3) insuring rapid dissemination of the data to potential users through

personal contacts, publications and workshops.

Accomplishment of the first specific objective started with the

publication of the Survey of Low Temperature Materials for Magnetic

Fusion Energy in March 1977. By continuing interactions with all low

temperature design, construction and measurement programs, such as the

Large Coil Project, we keep in touch with new problems as they arise.

This year's contribution to accomplishment of the second objective is

described here in Table 1 in the form of an outline of the research

projects. The results are contained later in this report. The third

objective is satisfied in part by these annual reports and by the handbook

project as well as by the series of NBS-DOE Workshops on Materials at

Low Temperatures which are held each fall.

3
Table 1. Outline of the NBS/DOE program on materials studies for

magnetic fusion energy applications at low temperatures.

Program Area Organi zation Program Description

A. Alloys for Low Temperature Service

1. Fatigue Martin Marietta Low cycle, strain controlled fatigue

testing of 304LN stainless steel and

5083-0 aluminum at 295, 76 and 4 K. 1

2. Fracture Martin Marietta Fracture toughness, fatigue crack

growth rate and tensile tests at 4 K on 2219-T87 aluminum alloy.

NBS Fracture toughness, fatigue crack

growth rates and tensile tests at 295, ' 76 and 4 K on AISI 304, 316, 304LN,

316LN and Nitronic 50 alloys. A total

of seven alloys were tested.

3. Test variables at 4 K NBS Effects of strain rate and discontin- ' uous yielding on strength, temperature

rise and martensitic products in

selected stainless steels examined.

4. Effects of temperature NBS

on strength

Tensile tests of AISI 304L conducted

at some 20 different temperatures ' between 4 and 295 K and results related

to martensitic transformations

.

B. Physical Properties of Stainless Steels

1. Elastic properties NBS Measurement of the elastic properties

of AISI 304, 304L, and 304N stainless

steels to determine: wi thin-specimen

variability (anisotropy), heat-to-heat

variability, effects of cooling to 4 K.

2. Magnetic properties Battel le Determination of the effect of magnetic

field and temperature on the magnetization of austenitic stainless steels,

and measurement of their magnetic

susceptibilities.

NBS Evaluate the initial permeabil ity

,

susceptibility and magnetic remanence

of selected stainless steels and their

welds at 4 K.

3. Thermal properties Wichita State U/

NBS

Measure the low temperature specific

heat of selected stainless steels.

4
Program Area Orqani zation Program Description

1. Metallurgical studies

2. Toughness tests

1. Characterization

1. Handbook preparation

2. NBS/DOE workshop

C. Stainless Steel Welding

Colo School of Characterize the microstructure, ferMines rite content and degree of sensitization of the 31 6L welds prepared by

Arcos

.

NBS Evaluate the fracture toughness of

31 6L as a function of ferrite content

at 295, 76 and 4 K. Conduct tests on 308L and 31 6L weld specimens supplied

by General Electric as part of their

large coil program. Conduct tests at

295, 76 and 4 K on AWS 330 welds supplied by Teledyne-McKay

.

. Nonmetal lies for Magnet Stuctures

NBS Characterization of the low-temperature mechanical , electrical, thermal

and elastic properties of G-IOCR and

G-llCR industrial fiberglass-epoxy

laminates including assessment of property variability among suppliers.

Cooperation with National Laboratories

in assessing radiation resistance of the

CR grades. Subsequent studies will

characterize less expensive NEMA laminates of the C and LE grades. A series

of unreinforced nonmetal lie materials

will be selected for detailed cryogenic

characteri zation

.

E. Technology Transfer

MPDC/NBS Production of a handbook containing

low temperature mechanical and physical

properties of materials of interest to

the MFE community. Presently designated for inclusion: AISI 304, 316

and 21/6/9 stainless steels; 2219,

5083 and 6061 aluminum alloys.

NBS An annual workshop to present research

results to the fusion community, to

discuss new problems and to promote

interaction between interested

parti es

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Highlights of Results

HIGHLIGHTS OF RESULTS

We present here a brief description of outstanding results from the

project reports which follow in the next section.

Alloys for Low Temperature Service

Emphasis was placed on evaluation of the mechanical properties of

ni trogen-strengthened stainless steels. Strain-controlled fatigue

performance of AISI 304LN compares favorably with AISI 304 and 316

steels, especially in the high cycle range at low temperatures . Fracture

toughness of AISI 304L and 316LN steels at 4 K remains high. The principal

advantage of the nitrogen-strengthened austenitic steels is the increased

yield strength at 4 K. Data were obtained documenting the effect of

nitrogen additions on yield strength at 4 K. The dependence of flow

strength on nitrogen concentration is stronger at lower temperatures.

Practical strain rates were shown to have little effect on 4 K

tensile yield strengths. The discontinuous yielding phenomena observed

in 4 K stress-strain behavior was confirmed to result from adiabatic

heating of the specimens.

Physical Properties of Stainless Steels

Elastic-constant variability in stainless steel 304 alloys, including

304L and 304N, was determined experimentally using an ultrasonic pulseecho-overlap method, which has an estimated uncertainty of + 0.1 percent

for these materials. Three kinds of variations -- sample-to-sample,

directional within sample, and repeated measurements on a single sample --

were determined for four elastic constants; bulk modulus. Young's

modulus, shear modulus, and Poisson's ratio. Based on 20 samples acquired

randomly, the overall variabilities for these four elastic constants

were 1.4, 0.9, 1.1, and 1.4 percent, respectively. Directional variations

within a sample never exceeded 0.5 percent.

9
Upon cooling below room temperature, the elastic-stiffness constants

increased regularly to about 40 K, where they decreased due to a Neel

(paramagnetic-antiferromagnetic) transition. Poisson's ratio shows

opposite behavior, first decreasing smoothly and then increasing near

40 K. Cooling curves for all studied samples were qualitatively similar,

but showed measurable differences both in the high-temperature slope

dC/dT and in the elastic-constant change AC due to the magnetic transition.

The bulk modulus (reciprocal compressibility) changes much less with

temperature than do the shear-type elastic constants.

Measurement of the magnetization and susceptibility of a large

number of stainless steels as a function of temperature and magnetic

field has led to a new interpretation of the magnetic state of these

materials at low temperatures. They are apparently all spin-glass

metals. In most cases the magnetization is not linear with field, which

leads to some ambiguity in defining the permeability and, in some instances,

significant amounts of hard ferromagnetism are found. Preliminary

magnetic properties measurements have been made on a series of 316 and

31 6L weldments containing varying amounts of ferrite. The remanence and

susceptibility values correlate well with magnegage measurements in most

instances but some significant deviations have been observed probably

resulting from a nonuniform ferrite distribution in the test specimens.

The specific heat of AISI 304 and 304L stainless steel over the

temperature range 2-20 K was measured for the first time. The results

indicate that, as expected, the specific heat is not greatly different

from that observed for the other austenitic stainless steels.

Stainless Steel Welding

Research continued on the investigation of metallurgical factors

contributing to the toughness (and degradation therof) of austenitic

10
stainless steel weld metals. Fracture toughness and metallurgical

studies on the series of 31 6L test welds prepared by Arcos Corp. have

been completed. The fracture toughness at 4 K ranged from 98 Mpav^ (90

ksi/Tn) for the welds with 10% ferrite to 179 MPavii? (162 ksi/fn) for the

ferrite-free welds; the toughness decrease was approximately linear with

ferrite content between zero and 10%. Metallurgical studies were conducted

by Colorado School of Mines and included microstructural characterization

,

ferrite content determination and sensitization studies. The fracture

toughness evaluation of AWS 330 welds prepared by Tel edyne-McKay at 295,

76 and 4 K was completed. The toughness at 4 K was 143 MPavfiT (130

ksi/Tn”). Tests were conducted on shielded metal arc deposits of 308L

and 316L weld metals using specimens provided by General Electric as

part of their involvement in the Large Coil Project. The results were

confusing in that the 308L weld had marginal Charpy impact toughness at

76 K (30 J; 22 ft lbs) yet the fracture toughness at 4 K was relatively

high (156 MPa/iT; 142 ksi/in). Furthermore, the Charpy toughness at 76 K

of the 316L welds from GE was higher than the values for the 31 6L Arcos

welds at the equivalent (zero) ferrite level, yet the fracture toughness

at 4 K was significantly lower at 95 and 118 MPa/iT (86 and 107 ksi/TrT)

for the GE welds and 178 MPa/TT (162 ksi/TrT) for the Arcos welds.

Further work is needed to identify the cause of the reduced toughness.

Nonmetal lies for Magnet Structures

A cooperative program between NBS and the U.S. laminating industry

has been established to provide commercial availability of special

grades of glass-epoxy, high-pressure industrial laminates intended for

use in cryogenic magnet structures. These products, identified as GlOCR and G-llCR meet present NEMA and military specifications for G-10

11
and G-11 laminates. The essential difference is that the CR grades are

produced to a common component and process specification so that property

variability is minimized.

The specifications for producing the CR grade laminates have been provided to all U. S. laminating firms through the Industrial Laminates Subcommittee of the National Electrical Manufacturers Association. The CR grades are

presently in commercial production. Production specifications will remain

proprietary within the U. S. laminating industry. Inquiries about the

specifications should be directed to Mr. H. Colin Smith, NEMA, Washington,

D.C. (202-457-8419).

The mechanical, thermal and electrical properties of one lot of GlOCR and G-llCR have been measured at 295 K, 76 K and 4 K by NBS and by

NBS subcontractors . The performance was found to equal that of the NEMA

G-10 and G-11 products. The characterization program on the CR grades

will continue in FY-79, including assessment to variability between

manufacturers . The characterization will include studies to assess the

benefits of fabricating the CR products with boron-free glass reinforcement for use in a neutron/gamma radiation environment.

Technology Transfer

Our major effort in this area is production of a handbook of low

temperature data on the mechanical and physical properties of selected

materials. The handbook is now in production, covers have been received,

and a first chapter (AISI 316) draft has been evaluated.

Our annual NBS/DOE Workshop on Materials at Low Temperatures was

attended by 70 people representing the materials and magnet design

12
communities. Indications are that the meeting was very productive for

those concerned.

The preparation of papers reviewing the state of the art in low

temperature materials selection and application is another aspect of

technology transfer. One such paper, discussing a wide range of structural

alloys, is included in this report.

13

Alloys for Low Temperature Service
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TESTING OF 2219-T87 ALUMINUM ALLOY AT 4 K

Martin Marietta

17

TESING OF 2219-T87

ALUMINUM ALLOY AT 4 K

J. A. Shepic

Principal Investigator

C. F. Fiftal

Program Manager

MARTIN MARIETTA CORPORATION

Denver Division

Post Office Box 179

Denver, Colorado 80201

19
INTRODUCTION

Aluminum alloys have been proposed for structural applications in MFE

magnet systems. One of the preferred alloys for cryogenic use is 2219 in

either T87 or T851 condition, which represent solution annealed, cold

worked to different strain values and aged. However, there is presently a lack of sufficient tensile and fracture data at 4K on this alloy,

particularly for sections greater than 1 inch. The objective of this work

was to characterize the tensile and fracture properties of heavy section

(1.5 inches thick) 2219-T87 plate at 4K. Transverse and longitudinal crack

growth parameters were also determined.

MATERIAL

The 1.5 inch thick 2219-T87 plate was procured from Martin Marietta

inventory. Quality Control inspection of the material established 295K

yield and ultimate at 55.1 ksi and 68.6 ksi, respectively, for the longitudinal grain direction. Elongation measured over a 2-inch gage length

was 8%. Chemical analysis is as follows:

El ement Mn Cu Zn Fe Ti

Specification 0.20-0.40 5. 8-6. 8 0.10 max 0.30 max 0.02-0.10

wt. %

Martin Marietta 0.20 6.38 0.008 0.11 0.034

analysis wt. %

Hardness (average of 5 readings) is 78.4 Rg.

MECHANICAL TESTING

Tensile specimens were taken from L, T and ST orientations. The L and T specimens were of standard threaded-end configuration, utilizing a 0.31 inch gage diameter and a 1.5 inch gage length. The ST tensile specimens were miniatures, 1.5 inches in overall length with a gage length and

diameter of 0.75 and 0.188 inches, respectively. The larger specimens

were instrumented with nichrome element strain gages while a special re- motely detachable 0.5 inch gage length extensometer was utilized on the

20
miniature specimens. Testing was performed in the cryostat and load

frame assembly pictured in Figure 1. Tensile data at 4K are summarized

in Table 1

.

The testing, 1.25-inch thick compact tension specimens were machined in the LT AND TL orientations. A WOL specimen configuration

shown in Figure 2 was chosen for the SL orientation since plate thickness was a limiting factor in specimen design. The WOL specimen B dimension was nominally 0.75 inches. Precracking was done at 295K at a low stress intensity level in accordance with ASTM E399-74. The maximum

stress intensity at the conclusion of precracking was less than 10 ksi/Tn

(11 MPa»/m) crack front curvature on all specimens was within the specification outlined in ASTri E399-74. Fracture testing of the compact tension

specimens was performed in a 100 kip MTS testing machine equipped with a liquid nitrogen jacketed vacuum dewar (Figures 3, 4). The WOL specimens were tested in the small glass dewar system used for the previously

described tensile tests.

Table 2 lists the results of the tests. All specimens satisfied the validity criteria for testing as specified in ASTM E399-74.

Fatigue crack growth rates were determined on TL and LT orientations utilizing 1-inch thick compact tension specimens. Testing was accomplished in the small test system since it allows easier access to

the specimen and is more efficient in the use of He. A compliance gage

was attached to integral knife edges machined in the face of the specimen

to facilitate crack growth detection at 4K. The cyclic frequency was

9 Hz.

Figures 5 and 6 are plots of da/dN vs AK, the fatigue-crack growth

data, which is summarized in Tables 3 and 4.

21
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2‘j

AK

calculation

is

based

on

95%

ofP

since

a

P

.

of

0.05

P_

was

used

throughout

the

test.

max

mm

max
FIGURE 1 System for Tensile and Crack Growth Testing at 4K

26
p

k

B =

3o

2H= 2. OB

D = .6B

For a/W from 0. 25 to 0. 65

1/2

W

39. 70 - 294. 2 e) + W

2 3 1118 (i) - 1M2 (^» + 1159

Notch Details

//-.003"Max. R

U

2

JB

Fatigue Crack Extension, Min

Figure 2 WOL Specimen Configuration

27

OO

loo
FIGURE 3 Overall View of Liquid Helium Testing System

20
FIGURE 4 Close-Up View of Load Train for Liquid Helium Testing System

20
5(MPa/m) 20

. I

1 - —

1-.-

i - -i- --T --

I

! 1

I....

~|

mm

cycle

i

i

!

'

i L • ^ 15 20

AK (ksi y/in)

Figure 5 da/dN vs. TL Orientation

Temp : 4K

25 30

30
da/dN

(n-in/cycle)

31

LOW CYCLE FATIGUE TESTING OF 304LN STAINLESS

STEEL AND 5083-0 ALUMINUM ALLOYS

Martin Marietta

33

LOW CYCLE FATIGUE TESTING OF 304LN

STAINLESS AND 5083-0 ALUMINUM

ALLOYS

J. A. Shepic

Principal Investigator

H. J. Brown

Program Manager

MARTIN MARIETTA CORPORATION

Denver Division

P.O. Box 179

Denver, Colorado

35
FOREWORD

This report de

Denver Division ot

November unde

scribes technical activities conducted by the

Martin Marietta from 26 July 1978 through 26

r National bureau of Standards (NBS) contract

CST-845b.

'Ine views and conclusions

of ttie authors and should not

senting the orficial policies

United States Government.

contained in this

be interpreted as

either expressed

document are those

necessarily repreor implied, of the

36
INTKODUCTION

Superconducting machinery will be subjected to cyclic loading

at temperatures down to 4 K. Fatigue data for candidate materials

are sparse oelow 20 K. At 4 K, most fatigue data are embodied in

tour papers (.references 1, 2, 3, and 4).

Stainless alloy AlSi 304LN and aluminum alloy 5083-0 were

evaluated. No previous fatigue data for either alloy at 4 K were

available.

MATERIALS

Eacn alloy was supplied by NBS as plate stock. The 304LN was

ootained as 3.8 cm (1.5 in.) thick plate. The 5083-0 aluminum

alloy was supplied in 2.5 cm (1.0 in.) thickness.

Cneraistry atid mechanical properties for 304LN, heat number

253^7, are given below:

L = 0.024 w/o

iTn = 1.74 w/o

P = O.U3o w/o

S = U.OUo w/o

Si = 0.51 w/o

Ni = 8.bU w/o

Cr = 18.05 w/o

Mo = O.JO w/o

Co = 0.23 w/o

Cu = 0 . JU w/ u

N = 0.13 w/o

Fe = ba lance

37
293

Temperature, K

77 4

Ultimate strength,

Hh/Tu2 i^ksi;

617 (90) 1541 (223) 1686 (245)

Yieia strengtn, 0.2% offset

diN/m2 i^KSi)

283 (4U 655 (95) 771 (112)

Elongation, % 63 47 37

Reduction of area, % 84 58 52

38
T

E

STING TECHNIQUE AND APPARATUS

Ttiis cnapter aescribes the specimen geometry, testing apparatus, and tecnnique used to obtain the fatigue data presented in

this report.

The fatigue specimen configuration used for this program is

similar to tnat used by Martin Marietta in previous fatigue studies freterences i, 2 and ftie design incorporates a smooth

gage section witn a mild radius and a bearing shoulder at the end

ot tne threaded section for alignment and locking, as opposed to

aligning on tne threads. This system is advantageous because only

two close-tolerance alignment surfaces are required and failure

witnin the prescribed area of the gage is assured. When locked

with a nut at each end of the threaded bar, the specimen can be

subjected to a fully reversed loading cycle that is continuous

tnrougn zero load. Figure 1 gives the specifications for the

spec imen.

Fatigue specimens were polished after machining using 240 to

J2U grit polishing paper (.the coarseness of tne initial rough polisaing paper depended on tne alloy haruness and machined surface

roughness in the gage section.) and then using 320 and 600 grit

paper followed oy a felt wheel and Tripoli polisning compound.

The polishing paper was bonded to a rubber sanding disc mounted in

a small Dumore grinding tool.

As tlie specimen is rotated at approximately 350 rpm between

lathe centers, the periphery of the spinning (approximately 20,000

rpm) sanding disc contacts tne specimen gage section with axial

39
I

40

Specifications

for

Fatigue

Specimen
s CroK.es

.

Cine same procedure was used with all grits). The final

polish uses a 2 in. diameter felt wheel mounted in a Dumore grinding cool ana loaded with Tripoli polishing compound. As the specimen is rotated between lathe centers at approximately 150 rpm,

tne spinning (approximately 20,000 rpm) impregnated felt wheel

contacts tiie specimen gage section with axial strokes. This polisuing Ceciinique produces a mirror finish in which any remaining

polisning marks are in an axial direction tiiat will have minimal

eftect on axial fatigue test results.

Because of the nonuniform cross section area of Che specimen,

It IS necessary to control total strain based on diametral deflection rather than axial deflection. A diametral strain exCensometer was aesigned and built specially for this work. This device

IS a one-piece aluminum ring with integral strain beams instrumented witn nichrome resistance strain gauges. The beam ends are

machined with knife edges to grip the specimen. The knife edge

separation distance is controlled to provide good gripping action

on the gage section. Tne gauge grips the specimen tight enough

tiiat it remains in tne liorizontal plane during filling of the

cryostat and subsequent cycling loading. Figure 2 gives a closeup

View of tne extensometer

.

Tne test configuration for 293 K evaluation was conventional.

AS stiown in Figure 3, tne specimen is installed in adapters that

permit the snoufder to be preloadea. Adapters are then installed

in tne testing machine.

41
Figuve 2 D-iametraZ Strain Extensometer

42
t'-igure 3 Apparatus for 293 K Testing

43
Testing at 77 K and 4 K was performed in a cryostat system

prepared tor iow-cycie fatigue testing. Loading was achieved

through two concentric frames shown in Figure 4. The smaller

frame surrounding tne specimen is constructed from titanium

alloy. Note that a specimen adapter plate is used at the top of

the specimen. Ht the bottom, the specimen shoulder bears against

the triangular plate. The remainder of the frame is constructed

trorn aluminum alloy. A glass vacuum dewar is used to provide environmental control. Figure 5 shows the load frame installed in

the machine, botli with and without the dewar in place.

The liquid nelium system is designed so testing can be

achieved witii minimum cryogen consumption. Tl»e titanium frame and

specimen are first precooled to 77 K and placed in a liquid nitrogen liilecl dewar on tne bottom platen. After cooling, the head is

raised and a liquid helium filled dewar replaces the liquid nitrogen dewar. A liquid level sensor is used to monitor level, and

refilling is performed as required.

Fatigue testing was performed in a 222 kN (50 kip) closed-loop

servo-controlled testing machine. Operation was in the straincontrol mode. The re lat lonsiiip between total axial strain amplitude and total diametral strain amplitude is as follows:

AP

= 2AX^ + ~ (1-2 m)

wnere ~ total axial or longitudinal strain amplitude

= total diametral strain amplitude

AP = load amplitude at half-life

44
Figure 4 Load Train for 4K Testing

45
i 4G

Figure 5

Apparatus

for 4K

Testing
A = cross-sectional area

E

= Young '

s

modulus

y = Poisson

s ratio

To ca Icu

1

ate the diametral strain amp 1 i tu de required

achieve a de s ired axi al strain amplitude, an approximate

XP was selected. Foi lowing completion of the test, a mo

ate val ue of AP was s elected from the load vs strain hys

curves and the actual total axi al stra in ampl itude value

calcula ted. Accuracy of the AP e s t ima tion im proved with

ceed mg test and, as a result. ac tua

1

values were close

values. Tne elast ic component of the total s train was c

using t ne AP value re ad from the test record and the app

value o f Young's modu lus. The plastic component of stra

calcula ted Dy subtree ting the eilast ic strain from the to

to

value for

re accurteres is

was reeach sucto desired

alculated

ropr iate

in was

tal strain.

EXPERIMENTAL TEST RESULTS

FatLgue benavLor of each of the two alloys was characterized

at 2^3, 77, and 4 K.. Because of the extreme high cost of liquid

tielLum and the large quantities required to obtain long-time fatigue data, only a limited number of 4 K tests were conducted beyond

10^ cycles.

304LN Stainless Steel - This alloy exhibits a significant

amount of discontinuous yielding at 4 K, as shown by prior tensile

test investigations. During cyclic loading, it is possible that

ttie heat generateu in the plastic range can cause vaporization of

47
the low heat capacity liquid helium and a resultant elevation of

the test temperature. This situation should be rate sensitive

and coulu be reduced by cycling at slow rates. Rate sensitivity

measurements made by Martin Marietta in previous work on 304

stainless (reference 11 show a frequency dependency above 0.1 Hz

at nigh cyclic strain amplitudes (1.3%1. Cyclic frequencies on

the bU4LiN latigue tests discussed here were held at or below 0.1

tiz tor strain amplitudes exceeding i.0% in liquid helium.

Experimental data for the 304LN alloy exhibited a temperature

dependency of Lite in the low cycle range in which life decreased

witn decreasing temperature. Below strain amplitudes of approximately 0.8%, life increased with decreasing temperature. Figure 6

grapnically depicts the fatigue life as a function of strain.

Table 1 gives detailed information on the diametral and longitudinal strains, stress amplitude, cyclic rate, and cycles to failure.

Figure 7 illustrates typical cyclic behavior of 304LN at room

temperature and a strain amplitude of 1.52%, fully reversed (R =

—1 ; . Loading was initiated in compression to reduce the possibility ot developing a local necked region during the first tensile

loading. Materials that strain soften are especially susceptible

to necking if tensile loading is appplied initially. Diametral

deflection is plotted on the x-axis vs load on the y-axis. The

calibrations lor each are shown on the graph. The first cycle

starts in compression at the center of the hysteresis loop. Note

tnat the load increases with each successive cycle as a result

ot strain naruening. The degree of strain nardening is dependent

on strain amplitude and test temperture.

48
O

onO

O

o

% ‘ apn:iT-[dniv pBU-tpnn-tSuoT I^^ox

w

a) r—

(

o

>. u

(1) 4^

'H

hJ

(U3

W)

'H

iJ b

49

Figure

6

Fatigue

Life

Curves

for

304LN
Table

1.

Fatigue

Data

for

304LN

Stainless

Steel \

50

^Specimen

Buckled

tTest

Discontinued

Temperature,

K.

293

77

‘V

Moauius,

GN/m2

psi)

195

(28)

209

(30)

207

(29.7)

Possion's

Ratio

0.290

0.279

0.271
Figure 7 Load-Defteotion Curve for Specimen 304 LN-?

51
Figure 8 illustrates a strain softening effect. Note that the

load excursion is lower at 48U0 cycles than at the start of the

test. The strain amplitude was 0.51% and the test was conducted

at room temperature.

A marked change in behavior occurs when the test temperature

IS reduced. Figure y shows a hysteresis plot of a 304LN specimen

testea at 77 K at a strain amplitude of 1.57%; the same strain

amplitude as Figure 7. The high degree of strain hardening is the

result of strain - induced martensitic transformation.

Results obtained at 4 K. for 304LN exhibit discontinuous yielding Dehavior. Figure 10 illustrates the first two cycles of a

test at 4 and a strain amplitude of 2.0%. The discontinuous

yielding is pronounced . Note that load drops occur only after a

small amount of initial yielding. This discontinuous behavior is

a result of adiabatic heating of the specimen due to straining,

causing a loss of strength.

Figure 11 illustrates benavior at a lower strain amplitude.

The load drops are not as severe and control of the test is improved. After 100 cycles nave elapsed, the hysteresis curve is

smooth

.

Figures 12 and 13 illustrate test results of even lower strain

amplitudes. Strain hardening is still taking place, however, discontinuous yielding is not apparent at the 0.5% strain level (Figure 137.

5063-0 Aluminum Alloy

Figure 14 summarizes the strain vs life data at the three

52

Figure 9 Load-Deflection Cierve for S-pecimen 304LN-8

54
Figure 10 Load-Deflection Curve for Speeinen 304LN-18 (First 2 Cycles)

b5
Figure 11 Load-Befteotion Curve for Speaimen 304LN-14

56
F-igure 12 Load-Deftecti-on Curve for Specimen 304Ld-16

295

' 1

59
temperatures. The trend is similar to the tested stainless alloys

with a decrease in cyclic lire at 4 K and hign strain amplitudes

witn a reversal of this oehavior at strain amplitudes below 0.9%.

The 293 K and 77 K. data at strain amplitudes above 1.5% are similar. These data are given in Table 2.

Figures 15 and lb are typical hysteresis curves for this alloy

at ambient temperature, strain softening was not apparent at any

or the strain amplitudes testea.

Figures 17 and 18 are indicative of cyclic behavior at 77 K.

Tests conducted at 4 K exhibit extensive discontinuous yielding (Figure 19). Changes in cyclic frequency by a factor of 5 between specimens 21 and 33 (that were tested at a strain amplitude

of 1.11%) did not show a significant degree of frequency depenaency. The same was done witn specimens 23 (Figure 20) and 24 at

0.75% strain amplitude with no significant difference in cycles to

raiiure. it can be reasoned that the aluminum alloy, because of

Its higher tnermal conductivity, will not be as susceptible to

cyclic trequency as the stainless.

Failure modes of the aluminum alloy are interesting. The

specimens fail in shear at strain amplitudes above 1% at 293 K and

77 K.. Failure is abrupt, with load decay occurring one or two

cycles before failure. Slight evidence of fatigue cracking is apparent on visual examination of the fracture surface.

Tne 4 K specimens that were tested at strain amplitudes above

1%, exhibit de lamination, much like a high strength, precipation

hardened aluminum alloy tensile tested at 4 K or 20 K.

60
Table

2.

Fatigue

Data

for

508Z-0

Aluminum

Alloy

Cl

Temperature, K

293

774

Moaulus,

GN/m2

pax;

71.6

(iO.4)

80

Ul.b)

81

(11

Possion's

Ratio

0.333

0.320

0.318
Fiyure 15 Load-Deflection Cupve for Specimen 5083-8
63
64
6b
G6
Figure 20 Load-Deflection Curve for Specimen 5083-23

67
At all temperatures and strain amplitudes near 0.6% or less,

tatigue cracking normal to the axis of loading takes place, followed by failure of the remaining ligament.

Discussion of Results

Data obtained trom this work were compared with other available information to obtain a comparison with alternate candidate

materials

.

For the jU4liN alloy, tne strain controlled data obtained from

this work were plotted with strain controlled data obtained for

304, 31b, ana 21-b-4 stainless alloys in a previous Martin Marietta test program ^Reference 1). A comparison of the fatigue

strength of 304LN with the other three alloys was performed at 293

R (Figure 21). 304DN is similar in behavior to the 21-6-9 alloy.

3U4LN nas better fatigue life than 304 at strain amplitudes above

1.0%. At 77 K (Figure 22), 304LN has lower fatigue life than 316

or 21-6-9 at strain amplitudes above 1.0% but is still better than

304. At higher cyclic lives (above 10,000 cycles), 304LN is superior to the other three alloys. Above 0.8% strain amplitude at 4

K. (Figure 23), the cyclic life of 304LN is similar to that of 304

and 2l-b-9, witn 31b superior in this range. However, at strain

amplitudes below 0.8% (cyclic life greater than 10,000 cycles),

304LH IS superior to 304 ana 21-6-9. Insufficient data on alloy

31b at nigh cyclic life prevent a comparison.

Figure 24 illustrates the strain nardening or strain softening

response as a function of temperature and strain amplitude for the

stainless alloys discussed above.

60
G9
% ‘apn^jxdiiiv uxeaxs x^^°I

70

Figure

22

Comparison

of

Fatigue

Data

at

77K
71
304; 4 K

600

0 1.0 2.0 3.0 Total Axial Strain, percent

Figure 24 Effect of Strain Amplitude on Change in Stress Amplitude

72

Stress

An^)litude,

Ao,

ksi
The 3083-0 fatigue data generated from this work were compared

with strain-contro i Led tatigue data determined lor aluminum alloys

2014-Tb and 22i!si-T851 by Nachtigall in ly?4 (Reference 3). This

comparison is shown grapnically at each test temperature in Figures 23, 2b, and 27. In all cases, at the higher strain levels

(low cyclic lives), 5083-0 has higher fatigue life than 2014-T6

but IS lower than 2219-1851. At low strain levels (high cyclic

lives), the alloys are comparable, with 2014-To having a slight

advantage in fatigue life.

Fatigue behavior of the 304LN and 5083-0 alloys do not display

qualities that should affect their candidacy for service at 4 K.

73
2219-T851

Figure

25

Comparison

of

Fatigue

Data

at

293 K
% ‘apn:jTxdtiiv ujpaqs TBUxpn^jSuoT Tb^ox

w

(U rHou

0) iw

•H

iJ

<u3

00

•H

4J

nJ

75

Figure

26

Comparison

of

Fatigue

Data

at

77K
o

t-H

w

0) rHo

CJ

CU

M-l •H

0)

:3bO

•H

4-) rt

o

o

% ‘apn^TjcIuiv UTBa^s IBU-rpn^TSuo^ Tb:iox

7G

F'Lgure

27

Compccrt-son

of

Fatigue

Data

at

4

K
KtFEkENCES

1. F. R. Schwarczberg and J. A. Shepic, "Fatigue Testing of

Stainless Steeis", Materials Studies for Magnetic Fusion

Energy Applications at Low Temper tures , Volume I, April 1978,

National bureau of Standards, Boulder, Colorado.

2. F. R. Scnwartzberg and T. F. Kiefer, "Study of Fracture Beliavior or Metals tor Superconducting Applications", Materials Research for Superconducting Machinery, Semi-Annual Technical

Report IV, October 1975, National Bureau of Standards,

Bouluer, Colorado.

d. A. j. Nacntigall, "Strain-Cycling Fatigue Behavior of Ten

Structural Metals Tested in Liquid Helium, in Liquid Nitrogen,

ana in Ainoient Air", NASA IN U-7532, February 1974.

4. A. J. Nacntigall, S. J. Klima, and J. C. Freche, "Fatigue of

oiquiu Rocket Engine Metals at Cryogenic Temperatures to

-4:;2F", NASA TN U-4274, December 19b7

5. E. J. becK. "Effect of beta Processing and Fabrication on Axial

Loauing Fatigue betiavior of Titanium", AFFDL-TR-69-108 , Martin

Marietta Corporation, June 1969.

7 /

FRACTURE AND STRENGTH PROPERTIES OF SELECTED

AUSTENITIC STAINLESS STEELS AT CRYOGENIC TEMPERATURES

National Bureau of Standards

79

FRACTURE AND STRENGTH PROPERTIES OF SELECTED AUSTENITIC STAINLESS STEELS

AT CRYOGENIC TEMPERATURES*

D. T. Read and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

*Work supported by DOE/OFE. Contributions of NBS, not subject to Copyright.

ol
ABSTRACT

Austenitic stainless steels have an excellent combination of mechanical and physical properties for load-bearing structures of large

superconducting magnets for plasma containment in magnetic fusion experiments. To assess their relative suitability, fracture toughness,

fatigue crack growth, and tensile properties data for five austenitic

stainless steels at 295, 76, and 4 K have been obtained. The steels

were AISI 304, 316, 304LN, and 316LN, and an Fe-21Cr-12Ni-5Mn alloy with

a higher nitrogen content than the other four grades. Tensile properties were measured at 4 K for two additional AISI 304L alloys with

higher than usual nitrogen contents. At 4 K the yield strength ranged

from 350 to 1330 MPa (49 to 188 Ksi). This wide spread in yield strength

was attributed to composition variations, especially nitrogen. A correspondingly large spread was observed in fracture toughness, from 131

to 328 MPa/m (119 to 298 ksi /in). The two principal findings were the

systematic variation of yield strength with nitrogen content and the

systematic inverse correlation between fracture toughness and yield

strength. Data from previous studies are reviewed which confirm the

trends of the present data. These results permit the selection of the

optimum combination of properties for a specific structural application.

82
INTRODUCTION

Nitrogen-strengthened austenitic stainless steels have attracted

interest as possible cryogenic structural materials because of their

superior strength [1-4]; but data on their toughness at cryogenic temperatures are generally lacking, except for a few studies [5-7]. The

availability of high-strength, high-toughness alloys for cryogenic

service would contribute to increased structural safety and economy in

devices such as magnetic-fusion-energy systems [8]. As cryogenic structures become larger and magnetic fields more intense, materials with

higher strengths than the widely used alloy AISI 304 are necessary. In

fact, nitrogen-strengthened AISI 304LN and 316LN have been selected for

respectively, the magnet cases for the Magnetic Fusion Test Facility at

Lawrence Livermore Laboratory and the toroidal field coil being built by

General Electric Company for the Large Coil Project. Clearly, development of a data base on the structural properties of nitrogen-strengthened

austenitic stainless steels is needed to support materials selection and

design for the large cryogenic magnet structures of the near future.

MATERIALS

All materials in this study are candidate materials for cryogenic

structures. Chemical compositions are given in Table 1. The AISI 304

and 316 are from the Oak Ridge National Laboratory reference heats whose

mechanical properties have been extensively characterized [9,10] at room

and higher temperatures for several product forms. The AISI 304L, 304LN

and 316LN are from commercial heats. The Fe-21Cr-12Ni-5Mn alloy (designated Nitronic 50 by the manufacturer**) was solutioned-annealed in the

laboratory before specimens were machined. Final heat-treatment parameters

**This tradename is used only for clarity. No approval or endorsement

of any commercial product by NBS is implied.

33
and measured grain sizes and hardnesses are given in Table 2. Representative microstructures are shown in Figs. 1-5. All microstructures

appear very similar. The grain size of the AISI 304 was large (grain

diameter 210 microns).

SPECIMENS

Fracture toughness and fatigue crack growth data were obtained

using compact specimens of the type shown in Fig. 6. All specimens were

2.54 cm (1 in) thick. Small, round tensile specimens oriented transverse

to the rolling direction were used. Their gage length was 2.54 cm

(1 in) and the diameter of the gage section was 0.4 cm (.15 in). The

orientations of fracture toughness (TL) and tensile specimens relative

to the original alloy plate are shown in Fig. 7.

TECHNIQUES

Toughness

All toughness data reported here were obtained using a variant of

the single specimen J-integral technique [11], with at least two complete

tests for each material and temperature. It has been shown [12,13]

that the critical stress intensity values obtained using the J-integral

technique (denoted Kj^(J)) are equal to or less than those obtained

using the conventional technique for critical plane strain stress intensity, called the test and described in ASTM Standard Test Method

E-399 [14]. The value of Kj^(J) is given by [15]:

Kj^J) = / EJj^/(l-v^) (1)

where E is Young's modulus, is the measured critical J-integral

value, and v is Poisson's ratio. The advantage of the J-integral

G4
technique is the comparatively small specimen size required compared to

1 5 the test. Specimen size requirements for the J-integral test are :

B, b, a, > 25 J/a ( 2 )

where B and a are as shown in Fig. 6, b = W-a, J is the J-integral

value, and a is the flow strength, here taken as the mean of the yield

and tensile strength. For the test, the requirement is^^:

B > 2.5

2

(3)

where is the yield strength.

For the AISI 304 alloy, the minimum specimen thickness required for

a valid J^^ test at 4 K was approximately 2.5 cm (1 in), while a valid

test required a thickness of about 220 cm (85 in). This is an

extreme case because this alloy has a very high toughness and a very low

yield strength. Nevertheless the thickness requirements for the

test are generally much less severe than for the test.

Despite the comparati vely small specimen thicknesses required for

J-integral testing, none of the room temperature tests proved valid

because of insufficient specimen thickness. Similarly, many of the

liquid nitrogen temperature tests were invalid. The requirement on

specimen thickness [15] may be recast as a requirement on the J values:

J < Ba/25 (4)
where B is the specimen thickness and a is the flow strength, the average

of the yield and ultimate strengths. For a given specimen thickness the

maximum valid J value may then be calculated. If the measured J value

turns out to be less than the maximum valid J value, then the test is

valid. If the measured J value is higher than the maximum valid value,

the test is invalid. The calculated maximum value from Eq. (4) then

becomes lower bound of the range of possible actual values. In the

Results section, any lower bound value calculated in this way will be

designated "Lower bound value calculated from the validity requirement

using the controlling specimen dimension." Note that a test must be

conducted for this conclusion; a test produces either a valid value

or a lower bound value.

A significant increase in the efficiency of J-integral testing

resulted from the use of the single-specimen method [12]. This technique allows a critical J value to be derived from a test using one

specimen, rather than at least four specimens as required in the

multiple-specimen technique [15]. Although each single specimen Jintegral test should be duplicated, a significant reduction in test time

and number of specimens is realized.

In the single-specimen J-integral technique the specimen is precracked in tension-tension fatigue to a predetermined crack length. The

specimen is loaded quasistatical ly (load line displacement rate 50 ym/s

or less) under displacement control to a point well within the elastic

region. The specimen is than partially unloaded, and the incremental

compliance curve is recorded with appropriate amplification and signal

conditioning. A signal proportional to the change in load is subtracted

(electronical ly in the present case) from the load-1 ine-displacement8G
change signal; the proportionality factor is adjusted so that the unloading compliance before crack extension is displayed as a vertical

line on the graph of load change against load line displacement change.

The usual graph of load against load line displacement is recorded on

another x-y plotter. The specimen is gradually extended, with several

stops for recording the unloading compliance. Typically, ten to twenty

unloading compliances are recorded. The magnitude of the decrease in

load during the unloadings is restricted to ten percent of the total

load. J-integral values are calculated in the usual way [16] from the

load vs. load-line-displacement record at each unloading point. Crack

extension values, Aa, are calculated from all unloading compliance

curves except the first one, which is used to evaluate the compliance at

Aa=0. The (Aa,J) data are then plotted, and is evaluated graphically as the intersection of a curve through the data with the theoretical blunting line given by J=2oAa where a is the arithmetic mean of the

yield strength and the ultimate tensile strength. Often the data followed the theoretical blunting line. Deviations were not systematic,

and were due to data scatter or experimental error. An example of the

J-Aa curves obtained is shown in Fig 8. The offset of the data from the

theoretical blunting line through Aa = 0 is artificial; for evaluation

of Jj^, a blunting line parallel to the theoretical line but passing

through the data is used. In this technique, scatter cannot be introduced into the J-Aa curve by specimen-to-specimen variability. Such

variability was an unknown factor in tests using the multiple specimen

technique.

The main disadvantage of the single-specimen method is that the Aa

values obtained from the unloading compliance are systematically lower

87
than those found by direct measurement of the fracture surface and are

scattered. The bias can be seen to have a small effect on the evaluation of typically 2 to 10%, although the slope of the J-Aa curve

will not be given accurately. The bias is believed to be caused by the

curved crack front which is formed during crack extension. The scatter

tends to be counteracted by the large number of data points available in

these tests. This scatter may be decreased in the future by more advanced test techniques, especially digital acquisition of the data.

Experimental Uncertainties in Toughness

Uncertainties in toughness data are of three kinds: first, those

associated with the quantities directly measured, such as load, crack

length, and displacement; second, those which arise from specimen to

specimen and heat to heat material variability; and third, those associated with the procedure used to infer toughness values from the

measured quantities. Errors of measurement are readily assessed by

repeated measurement, measurement of standard loads or dimensions, and

other usual calibration procedures. Errors associated with material

variability may be the most severe. Evaluation of material properties

for structural design requires careful consideration of uncertainties

due to material variability, through understanding of its causes and

extensive data on its magnitude. The consistency obtained by establishment of standard procedures for toughness tests reduces relative

measurement errors. Since the standard procedures for elastic-plastic

fracture toughness testing are still evolving, however, a limited discussion of the inherent measurement errors in these tests is helpful.

Measurement uncertainties in individual J values are considered to

be of the order of 5% or less, from accumulated uncertainties in load

88
cell and displacement gauge calibrations and measurements of the area

under the load-displacement record. Values of crack extension, Aa,

from partial unloading compliances for the single specimen J-integral

test averaged about 70% of the actual Aa value, with a standard deviation of around 0.7 mm. The bias is believed to contribute an uncertainty of 10% or less to the measurement. The scatter is compensated by taking many data points, and does not further increase the

uncertainty in Jj^. The cumulative uncertainty in values is estimated to be +15% or less, so the measurement uncertainty in values

inferred is 8% or less since is proportional to /J . Data for

comparison of the single specimen J-integral technique to the multiplespecimen technique were obtained at 76 and 4 K for AISI 304N. The 76

and 4 K single specimen values were 2 and 13%, respectively, lower

than the previously obtained values [7], in agreement with the +15%

cumulative error in Jj^ estimated. Errors in fracture toughness associated with material variability may also be large. Data reported

here are, in general, not applicable to materials with different chemical composition or thermochemical processing.

Fati que

Fatigue crack growth rates were also measured using compact specimens. The specimen is loaded repeatedly (at a frequency of 20 Hz) with

the minimum load one tenth of the maximum. The compliance, the reciprocal slope of a plot of load against load-line-displacement, is recorded periodically at intervals ranging from 100 to 20,000 cycles

depending on the rate of crack growth. The compliance data are converted to crack lengths using the formula [17]

89
a = W[-.9505 + .3447(EBC)^/^ + .5654(EBC)^'^^ - . 1 748(EBC)^/^] (5)

where C = displacement/load and E and B have been previously defined.

This equation is a least squares fit, developed in our laboratory, to

compliance data given in a proposed ASTM test method [18]. The crack

length is also used in calculating the applied stress intensity factor

range using the formula [14]

= (WPm1n>T[29-6-l85(^)+655.7(^)2-10l7(^)3.638.9(^)^] (6)

where , and P^. are the maximum and minimum loads, max min

Strain gage, clamp-on extensometers were used in tensile measurements. Yield strength values at 0.2% offset are reported. The flow

strength is taken as the average of the yield and ultimate strengths.

Uncertainties associated with inferring property values from measured quantities are not a problem in fatigue and tensile property

measurements, because standard test procedures are widely accepted.

However uncertainties of measurement and material variability remain.

Standard deviations of 20 to 40% from the average value were typical in

the fatigue crack growth data. Fatigue data are not considered to be as

sensitive to material variability as toughness data, and in fact data

from two or three specimens are usually consistent. Measurement uncertainties in tensile properties are estimated to be 3% or less; these

are due to uncertainties in specimen dimensions, load values as read

from analog recorder traces, and (for the yield strength) specimen

strains from the calibrated extensometer.

For all tests, temperature control was achieved by immersing the

specimen in liquid helium or liquid nitrogen at local atmospheric

90
pressure* or by testing in laboratory air. The apparatus used for these

tests has been described by Fowlkes and Tobler [19].

RESULTS AND DISCUSSION

Fracture toughness results are listed in Table 3 as critical plane

strain stress intensity values calculated from J-integral test results,

denoted using Eq. (1). The 4 K fracture toughness results for

AISI 304 and 316 are significantly lower than some reported earlier

[20], but these two stainless steels remain tougher at 4 K than any

other material. Two possible reasons for the discrepancy are heat to

heat material variability and changes in test procedures. The previously tested materials had the same orientation as the present study.

The AISI 304 in the present study had lower hardness and lower yield

strength at 295 and 76 K; the AISI 316 in the present study had lower

hardness and lower yield strength at 295 K, but higher yield strength at

76 and 4 K. A thorough assessment of the heat to heat J-integral variability between materials cannot be made because in the present test

program the Aa values obtained are rather small, 6 mm or less, while in

the previous test program only a few of the directly measured Aa values

were comparably small. Some overlap in Aa values between those of the

present and previous studies is present for both AISI 304 and 316, and

in both cases the J-integral values obtained in the present study were

significantly less than those obtained at comparable Aa levels in the

previous study. The regions of overlap were around Aa=6 mm, while the

critical J values were determined from Aa values of about 2 mm and less

in the present study. It is concluded that the discrepancy in results

*Typically .082 MPa. At this pressure, helium boils at 4 K and nitrogen

at 76 K.

91
between present and previous studies is at least partially due to material variability, and may also be due in part to the changes in the Jintegral test procedure which now tend to restrict attention to relatively small values of the crack extension.

Fracture toughness is plotted against yield strength for selected

stainless steels [7,20] at 4 K in Fig. 9. Both fracture toughness and

tensile data for this plot were obtained on specimens from the same

plate. The selected data represent wrought and annealed austenitic

stainless steels. No weldments nor sensitized nor otherwise improperly

heat-treated alloys have been included. Fig. 9 shows that there is a

linear relationship between yield strength and fracture toughness for

austenitic stainless steels at 4 K. All weldments fall well below the

trend line shown in Fig. 9.

The relationship between fracture toughness and yield strength has

been reported by other observers [21-25]. The trend lines reported by

Pellini are shown in Fig. 10 along with the same data shown in Fig. 9.

The agreement between the trends is quite good, and probably fortuitous,

especially when it is noted that Fellini's data are for body-centeredcubic materials at room temperature while the data of the present study

are for face-centered-cubic materials at the temperature of liquid

helium. No satisfactory explanation for the relationship between

strength and toughness shown in Fig. 10 has been offered. However this

trend may be significant if it represents, for 4 K use, the metal quality

attainable with current metallurgical techniques. Perhaps better properties will be attainable or perhaps the trend represents the limit of

what can be achieved in wrought, annealed stainless steels. Further

research is needed to determine whether and how better combinations of

strength and toughness can be obtained.

92
Fatigue crack growth results at 4 K are indicated schematically in

Fig. 11. The complete list of fatiguq crack growth results is given in

Table 4. The rates for AISI 304LN are higher than those for the other

materials. Comparison of AISI 304LN and 304N [7] rates with 304L and

304 rates at 4 K indicate that the nitrogen addition increases the

fatigue crack growth rates by a factor of 4 to 5. However, this effect

is not observed in AISI 316 as both 316LN and 316 have equivalent rates

at 4 K. All the fatigue crack growth rates observed in the present

study were within the band of previously observed results for austenitic

stainless steels [7].

Tensile test results are listed in Table 5. Temperature dependences of the yield and ultimate strengths are shown in Fig. 12. The

alloys that have higher 4 K yield strengths have stronger temperature

dependences. Elongations and reductions of area are plotted against

temperature in Fig. 13. The general tendency of these quantities to

decrease with decreasing temperature is evident.

The effect of nitrogen content on the 4 K yield strength is displayed in Fig. 14. Data are included from other studies [7,20,26]. At

higher temperatures the solid solution strengthening effect is not as

pronounced. A consistent increase in yield strength with increasing

nitrogen content is apparent [27].

93
SUMMARY AND CONCLUSIONS

1. Fracture toughness, yield and ultimate tensile strength, and

fatigue crack growth rates at 295, 76, and 4 K have been obtained

for five austenitic stainless steels. Tensile data at 4 K were

obtained for two additional stainless steels.

2. At 4 K, the yield strengths ranged from 350 to 1330 MPa (49 to

188 ksi), while the fracture toughness ranged from 138 to 321 MPa

v4tF (119 to 298 ksi /Tn )

.

3. At 4 K, fracture toughness decreased linearly with increasing yield

strength.

4. At 4 K, yield strength increases linearly with nitrogen content.

This dependency on nitrogen concentration becomes very significant

at low temperatures. It appears that austenitic stainless steels

can be "tailor-made" for 4 K applications by adjusting the nitrogen

concentration to produce a desired yield strength-fracture toughness

combination.

94
REFERENCES

1. D. C. Larbalestier and D. Evans, "High Strength Austenitic Stainless

Steels for Cryogenic Use," Sixth International Cryogenic Engineering

Conference, Grenoble, France (1976), IRC Science and Technology Press

Ltd., Surrey, England, pp. 345-347.

2. J. W. Montano, "Stress Corrosion Resistance and Cryogenic Temperature

Mechanical Behavior of 18-3Mn (Nitronic 33) Stainless Steel Parent and

Welded Material," NASA TM X-73309, Marshall Space Flight Center, AL (1976).

3. C. E. Spaeder, Jr., and W. F. Domis, "Cryogenic and Elevated-Temperature

Mechanical Properties of a High-Nitrogen Type 304 Stainless Steel,"

Report No. 44JD12-062-2, United States Steel Corporation, Pittsburgh, PA

(1969).

4. R. B. Gunia and G. R. Woodrow, "Nitrogen Improves Engineering Properties

of Chromium-Nickel Stainless Steels," J. of Materials, JMLSA, (June

1970), pp. 413-430.

5. R. L. Tobler and R. P. Reed, "Tensile and Fracture Behavior of a NitrogenStrengthened, Chromium-Nickel -Manganese Stainless Steel at Cryogenic

Temperatures," Symposium on Elastic-Plastic Fracture, American Society

for Testing and Materials, Atlanta, GA (Nov. 1977).

6. P. P. Dessau, "LN

2 and LH^ Fracture Toughness of Armco Alloy 22-13-5,"

Report No. TID/SNA-2083, Aerojet Nuclear Systems Co., Sacramento, CA (1971).

95
7. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and Tensile

Properties of Three Nitrogen-Strengthened Stainless Steels at Cryogenic

Temperatures," in Material s Studies for Magnetic Fusion Energy Appl i cations

at Low Temperatures-I , NBSIR 78-884 (1978), pp. 93-154.

8. R. P. Reed, F. R. Fickett, M. B Kasens, and H. I. McHenry, "Magnetic

Fusion Energy Low Temperature Materials Program: A Survey, " in Materials

Studies for Magnetic Fusion Energy Appl i cations at Low Temperatures-I ,

NBSIR 78-884 (1978), pp. 245-335.

9. R. W. Swindeman, W. J. McAfee, and V. K. Sikka, "Product Form Variability

in the Mechanical Behavior of Type 304 Stainless Steel at 593°C," in

Reproductibil ity and Accuracy of Mechanical Tests

,

ASTM STP 626, American

Society for Testing and Materials, Philadelphia, PA (1977), pp. 41-64.

10. V. K. Sikka, "Product Form Characterization of Reference Heat of Type 316

Stainless Steel, ORNL-5384, Oak Ridge National Laboratory, Oak Ridge, TN (1978).

11. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, "Single

Specimen Tests for D'etermindtion , " in Mechanics of Crack Growth, ASTM

STP 590, American Society for Testing and Materials, Philadelphia, PA (1976),

pp. 27-42.

12. W. A. Logsdon, "Elastic Plastic (Jj^) Fracture Toughness Values: Their

Experimental Determination and Comparison with Conventional Linear Elastic (kj^) Fracture Toughness Values for Five Materials," in Mechanics

of Crack Growth , ASTM STP 590, American Society for Testing and Materials,

Philadelphia, PA (1976), pp. 43-60.

9G
13. D. T. Read and R. P. Reed, "Effects of Specimen Thickness on Fracture

Toughness of an Aluminum Alloy," Intn. J. Fracture 1^ (1977), pp. 201-

213.

14. Standard Method of Test for Plane Strain Fracture Toughness of Metallic

Materials (Designation E399-74), 1974 Annual Book of ASTM Standards , Part

10 (1974) 432-451.

15. J. D. Landes and J. A. Begley, "Test Results form J-Integral Studies: An

Attempt to Establish a Testing Procedure," in Fracture Analysis , ASTM

STP 560, American Society for Testing and Materials, Philadelphia, PA (1974),

pp. 170-186.

16. J. G. Merkle and H. T. Corten, "A J-Integral Analysis for the Compact

Specimen, Considering Axial Force as Well as Bending Effects," J. Pressure Vessel Tech., Trans ASME, No 6 (1974), pp. 1-7.

17. H. I. McHenry, National Bureau of Standards, Boulder, CO, private communication.

18. Proposed Recommended Practice for R-Curve for R-Curve Determination,

1974 Annual Book of ASTM Standards , Part 10, American Society for Testing

and Materials, Philadelphia, PA (1974), pp. 669-683.

19. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a Ti6A1-4V Alloy at Liquid Helium Temperature," Eng. Frac. Mech. 8 (1976),

pp. 487-500.

97
20 . Materials Research for Superconducting Machinery (I-VI, edited by R. P.

Reed, A. F. Clark, E. C. van Reuth and H. M. Ledbetter. NTIS Order

Numbers AD780596, ADA004586, ADA012365, ADA019230, ADA030170, and ADA036919,

National Technical Information Service, Springfield, VA (1973-1977).

21. R. J. H. Wanhill, "Microstructural Influences on Fagigue and Fracture

Resistance in High Strength Structural Materials," Engr. Fracture Mech.

10, (1978), pp. 337-357.

22. J. P. Hirth and F. H. Froes, "Interactions Between Fracture Toughness and

Other Mechanical Properties in Titanium Alloys," Metall. Trans. 8A,

(1977), pp. 1165-1176.

23. G. T. Hahn and A. R. Rosenfield, "Metallurgical Factors Affecting Fracture Toughness of Aluminum Alloys," Metall. Trans. 6A (1975), pp. 653-670.

24. M. G. Stout and W. W. Gerberich, "Structure/Property/Continuum Synthesis

of Ductile Fracture in Inconel Alloy 718," Metall. Trans. ^ (1978),

pp. 649-658.

25. W. S. Pell ini, R. J. Goode, P. P. Puzak, E. A. Lange, and R. W. Huber,

"Review of Concepts and Status of Procedures for Fracture-Safe Design of

Complex Welded Structures Involving Metals of Low to Ultra-High Strength

Levels," U.S. Naval Research Laboratory Report 6300, Washington, D.C.

(1965).

98
26. R. Voyer and L. Weil, "Tensile and Creep Properties of a High Nitrogen

Content 18/10 (AISI 304L) Stainless Steel at Cryogenic Temperatures,"

Adv. Cryo. Eng. lj_(1965), pp. 447-452.

O 27. L.-A Norstrdm, "The Influence of Nitrogen and Grain Size on Yield Strength

in Type A.I.S.I 316L Austenitic Stainless Steel," Metal Science 1J_ (1977),

pp. 208-212.

99
LIST OF TABLES

Table 1

Table 2

Table 3

Table 4

Table 5

Chemical compositions of the materials used in the present study in

weight percent, as supplied by the manufacturers except for AISI 30^,

g supplied by ORNL analysis.

Hardnesses and grain sizes of the materials used in the present study.

Fracture toughness results obtained in the present study.

Paris law fatigue crack growth rate parameters for the annealed austenitic stainless steels. The fatigue crack growth rate da/dN is given

by da/dN = C(AK)'^, where AK is the stress intensity range.

Tensile properties at 295, 76, and 4 K for the annealed austenitic

stainless steels.

100
Table 1. Chemical compositions of the materials used in the present

study in weight percent, as supplied by the manufacturers

except for AISI 304, supplied by ORNL analysis.

Designation Cr Ni fin Mo P S Si C N Fe

A.I.S.I. 304* 18.44 9.72 1.28 0.32 0.030 0.016 0.48 0.051 0.031 bal

A.I.S.I. 304LN 18.85 8.60 1.74 0.38 0.38 0.003 0.51 0.2^ 0.13 bal

A.I.S.I. 316 17.25 13.48 1.86 2.34 0.02^ 0.019 0.58 0.057 0.030 bal

A.I.S.I. 316LN 17.40 13.90 1.58 2.50 0.021 0.021 0.A8 0.016 0.16 bal

Fe-21Cr- 12Ni-5Mn 21.15 12.37 4.96 2.17 0.026 0.015 0.^9 0.041 0.310 bal

A.I.S.I. 304L 18.54 9.27 1.78 n . a

.

0.019 0.012 0.50 0.021 0.075 bal

A.I.S.I. 304L 18.7 9.2 1 .74 0.38 0.029 0.006 0.37 0.021 0.096 bal

*Data from Ref. 9.

101
Table 2. Hardnesses and Grain Sizes of the Materials Studies

Designation

Hardness

(Rockwell B)

Grain Diameter"*"

(microns

)

ASTM Grain

Size Number

A.I.S.I. 304 75 210 1 .5

A.I.S.I. 304LN 79 57 5.5

A.I.S.I. 316 79 65 5.0

A.I.S.I. 316LN 83 105 3.5

Fe-25Cr-12Ni-5Mn* 93 160 2.4

"''Measured by the intercept method.

A non-standard heat treatment was used for this material. It was annealed for

1.5 hr at 1150°C and water quenched.

102
Table 3. Fracture Toughness

MPa/m (ksi/Tn).

Results Obtained in the Present Study i n

Temperature (K)

Material 295 76 4

A.I.S.I. 304 368? (335)

365? (332)

544^ (500)

393^ (358)

404^ (368)

399^ (363)

317

338

328

(288)

(308

(298)

264“^ (240)

A.I.S.I. 304LN >608? (553)

557^ (507)

310 (282)

376 (342)

343 (312)

213

195

221

210

(194)

(177)

(201)

(191)

284*^ (258)

A.I.S.I. 316 371? (338)

486? (442)

387^ (352)

351^ (319)

418^ (380)

436 (397)

402^ (366)

280

267

256

268

(255)

(243)

(233)

(244)

271^

A.I.S.I. 316LN 328^ (298)

278® (253)

286^ (260)

255 (232)

277 (252)

266 (242)

210

225 , 275^

203

(191)

(205)

(259)

(185)

Fe-21Cr-l 2Ni -5Mn

(Nitronic 50)

287? (352)

337® (307

268 (244)

270 (246)

255 (232)

264 (240)

114

147

131

(104)

(134)

(119)

321^ (292)

^Invalid result according to the thickness criterion for a valid J-integral

test.

toughness lower limit J-, . from J, . = Ba/25 where B = specimen ligament

and a = flow strength. ^

Marginally invalid result according to the thickness criterion for a valid

test. Ligament is within 10% of the required value.

^Result from multiple-specimen analysis using an insufficient number of

specimens

.

103
Table 4. Paris lav^ fatigue crack growth rate parameters C and n for

annealed austenitic stainless steels. The fatigue crack growth

rate da/dN in mm/cycle (in/cycle) is given by da/dN = C(AK)^,

where AK is the stress intensity in MPa^m (ksi/Tn). The C values enclosed in parentheses are to be used when AK and da/dN

are i n Engl i sh units.

Material Temperature C n

A.I.S.I. 304 304 6.58

76 9.51

4 3.26

X 10"

p (3.76

X 10"

n (5.83

X 10" ^ (1.84

X 10" b 3.64

X 10"

f) 4.34

X 10"") 3.49

A.I.S.I. 304LN 295

76

4

3.01 X 10“q (1

2.67 X 10"^. (1

1.69 X 10"'^ (1

69 X 10" b 3.17

45 X 10" b 3.17

02 X 10"" ) 4.25

A.I.S.I. 316 295

76

4

6.89 X 10

4.41 X 10

8.33 X 10

(4.44 X 10

(2.19 X 10

(4.33 X 10

-13

II^

4.84

2.25

2.71

A.I.S.I. 316LN 295

76

4

5.66 X 10

3.01 X 10

7.56 X 10

(3.17 X 10

(1.49 X 10

(4.15 X 10

-11

3.43

2.24

3.26

Fe-21Cr-12Ni-5Mn 295

(Nitronic 50) 76

4

3.46 X 10"

^

(2

1.05 X 10"b (6

1.94 X 10" (1

49 X 10"]b 5.91

73 X 10" J 4.79

32 X 10"'^) 5.33

104
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107
LIST OF FIGURES

Fig. 1. Compos i te 3-dimensi onal photomicrograph of the AISI 304 alloy.

Fig. 2. Composi te 3-dimensi onal photomicrograph of the AISI 304LN alloy

Fig. 3. Composi te 3-dimensional photomicrograph of the AISI 316 alloy.

Fig. 4. Composi te 3-dimensional photomicrograph of the AISI 316LN alloy

Fig. 5. Composi te

(rii tronic

3-dimensi onal

50) alloy.

photomicrograph of the Fe-21Cr-12Ni-5Mn

Fig. 6. Specimen geometry.

Fig. 7. Orientations of compact and tensile specimens.

Fig. 8 Typical J-Aa curve obtained using the single-specimen J-integral

method. Inferred Aa values are systematical ly lower than actual

values after accounting for the zero shift.

Fig. 0. Fracture toughness at 4 K of selected wrought and annealed austenitic

stainless steels plotted against their 4 K yield strengths. The line

is a linear least squares fit.

Fig. 10. Fracture toughness at 4 K of selected wrought and annealed austenitic

stainless steels plotted against their 4 K yield strengths with least

squares fit line and Fellini's data band centers (dashed lines) for

"old" and "new" ferritic steels at room temperature.

Fig. 11. Best-fit lines for 4 K fatigue crack growth rate data for selected

wrought, annealed stainless steels.

Fig. 12. Temperature dependences of the yield and ultimate strengths of the

annealed austenitic stainless steels.

Fig. 13. Temperature dependences of the elongation and reduction of area of

the annealed austenitic stainless steels.

Fig. 14. Yield strength at 4 K for annealed austenitic stainless steels plotted

against weight percent nitrogen. Behavior trends at 76 and 295 K are roughly indicated by the dashed lines.

108
AISI 304

L - Longitudinal or rolling direction

T ' Long transverse or width direction

S ' Short transverse or plate thickness direction

Fig. 1. Composite 3-dimensional photomicrograph of the AISI 304 alloy.

109
AISI 304 LN

L - Longitudinal or rolling direction

T - Long transverse or width direction

S ' Short transverse or plate thickness direction

Fig. 2. Composite 3-dimensional photomicrograph of the AISI 304LN alloy.

no
AISI 316

L - Longitudinal or rolling direction

T - Long transverse or width direction

S - Short transverse or plate thickness direction

3. Composite 3-dimensional photomicrograph of the AISI 316 alloy.

Ill
AISI 316 LN

4V>

L - Longitudinal or rolling direction

T “ Long transverse or width direction

S - Short transverse or plate thickness direction

Fig. 4. Composite 3-dimensional photomicrograph of the AISI 316LN alloy.

11 ^
Fe>21CM2Ni-5Mn

(Nitronic 50)

L - Longitudinal or rolling direction

T • Long transverse or width direction

S - Short transverse or plate thickness direction

Fig. 5. Composite 3-dimensional photomicrograph of the Fe-21Cr-12Ni-5Mn

(Nitronic 50) alloy.

113
Fig. 6. Specimen geometry.

114
I

Rolling Direction

Fig. 7. Orientations of compact and tensile specimens.

115
-INTEGRAL

(kJ/m

Fig. 8 Typical J-Aa curve obtained using the single-specimen J-integral

method. Inferred Aa values are systematically lower than actual

values after accounting for the zero shift.

116
Fig. 9. Fracture toughness at 4 K of selected wrought and annealed austenitic

stainless steels plotted against their 4 K yield strengths. The line

is a linear least squares fit.

117
FRACTURE

TOUGHNESS

(MPa/nT)

YIELD STRENGTH (ksi)

300

200

100

0

to

W)oo

CJ3 rao

UJ

Q£ZD

<_) C

Fig. 10. Fracture toughness at 4 K of selected wrought and annealed austenitic

stainless steels plotted against their 4 K yield strengths with least

squares fit line and Fellini's data band centers (dashed lines) for

"old" and "new" ferritic steels at room temperature.

118
0}

UJ

h-<

a:

o

a:cno

<t

ce:

C_)

STRESS INTENSITY RANGE (ksi-in^")

10 20 50 80

-3

-4

-5

-6

Fig. 11. Best-fit lines for 4 K fatigue crack growth rate data for selected

wrought, annealed stainless steels.

119

CRACK

GROWTH

RANGE

(in/cycle)
Fig. 12. Temperature dependences of the yield and ultimate strengths of the

annealed austenitic stainless steels.

120

Strength

(ksi)
Elongation

or

Reduction

of

Area

(percent)

Fig. 13. Temperature dependences of the elongation and reduction of area of

the annealed austenitic stainless steels.

121
YIELD

STRENGTH

(MPa)

Fi'i. 14. Yield strength at 4 K for annealed austenitic stainless steels plotted

against weight percent nitrogen. Behavior trends at 76 and 295 K are roughly indicated by the dashed lines.

1 o

L.

YIELD

STRENGTH

(ksi)
HEATING AND STRAIN RATE EFFECTS IN

AISI 304L STAINLESS STEEL AT 4 K

National 'Bureau of Standards

1^3
r
HEATING AND STRAIN RATE EFFECTS IN AISI 304L STAINLESS STEEL AT 4 K*

D. T. Read and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, CO 80303

*Work sponsored by DOE/MFE. Contribution of NBS, not subject to copyright.

12b
ABSTRACT

In structural alloy tensile specimens undergoing mechanical v/ork in

a liquid helium (4 K) bath, internal temperature increases are favored

by low specimen specific heat and thermal conductivity. Indeed, the

well-known phenomenon of serrated yielding is considered to be due to

adiabatic heating of the specimen. Besides being of scientific interest, the possible occurrence of such heating must be accounted for in

testing structural materials for liquid helium temperature applications.

Therefore, temperatures of small tensile specimens of AISI 304L stainless

steel were measured during straining at 4 K at several strain rates.

Stress-strain curves were obtained at three strain rates. Serrated

yielding was observed at all strain rates. Temperatures above 4 K were

detected only during and immediately following the individual load

drops. The maximum temperature rise detected was 52 K. No effect of

crosshead speed was detected. Examination of a specimen which was

deformed only until one load drop occurred showed that during the load

drop only a small portion of the specimen deformed. Analysis of the

relative number of occurrences of temperature rises of different magnitudes showed that only a small portion of the specimen was heated

during a given load drop. The calculated size of the heated portion of

the specimen agreed with the visually observed size of the severely

deformed portion. The value of the maximum temperature attained was

calculated from the load drop, volume of specimen material heated, and

the physical properties of the specimen material; the calculated value

agreed well with the measured value.

12G
INTRODUCTION

In structural alloy tensile specimens undergoing mechanical work in

a liquid helium (4 K) bath, internal temperature increases are favored

by low specimen specific heat and thermal conductivity. Indeed, the

well-known phenomenon of serrated yielding is considered to be due to

adiabatic heating [1,2] of the specimen. Besides scientific interest,

the possible occurrence of such heating must be accounted for in testing

structural materials for liquid helium temperature applications. Therefore, temperatures of small tensile specimens of AISI 304L stainless

steel were measured during straining at 4 K at several strain rates.

Stress-strain curves were obtained at three strain rates.

MATERIAL

The material used in this study was a 2.5 cm (1 inch) plate,

A.I.S.I. 304L stainless steel obtained commercially in the mill-annealed

condition. Its chemistry is shown in Table 1. Flat tensile specimens

in the longitudinal orientation with a gage section 2.5 mm (0.1 inch)

thick by 4.8 mm (.19 inch) wide by 30 mm (1.2 inch) long were used for

all tests. The nominal grain diameter was 70 ym and the hardness

Rockwell B-83.

TECHNIQUES AND APPARATUS

The mechanical property tests were carried out using a screw-driven

tensile machine with the capability for variable crosshead speeds,

adapted for low temperature testing as described by Reed [3]. The

strain rates are obtainable directly from the crosshead speeds and gage

length under the assumption of negligible work hardening and negligible

strain rate outside the reduced section. Table 2 gives the strain rates

which correspond to the crosshead speeds. When the work hardening

127
coefficient is nonzero, the actual strain rates are less than those

calculated from the crosshead speeds. During load drops, the actual

strain rates are quite rapid, and, as far as is known, are independent

of crosshead speed. Therefore, crosshead speeds will be given rather

than strain rates when describing experiments and data. Displacement

gages as described by Reed and Mikesell [4] were used to measure specimen strain.

For detection of temperature rises during deformation, the specimen

was instrumented as shown in Fig. 1. A NbTi superconducting wire with a

transition temperature of about 10 K was bonded into a slot milled into

the specimen. This allowed detection of the occurrence of temperature

rises along the whole length of the specimen, but provided no information on their magnitude. An Au-.05 at % Fe versus chromel thermocouple

was soldered to the specimen. This allowed measurement of the amplitude

of the temperature rise a specific location, and, as discussed below,

allowed calculation of the temperature profile of the load drops.

Electrical signals corresponding to the resistance rises in the superconducting wire or emf's from the thermocouple were plotted using an xy plotter with a response time of a few tenths of a second.

Martensitic transformation is associated with plastic deformation

at low temperatures in AISI 304L. Since the martensitic, body-centered

cubic phase is ferromagnetic , local regions of deformation along the

specimen were conviently decorated with iron powder.

RESULTS AND DISCUSSION

Serrated yielding occurred at all crosshead speeds for specimens

immersed in liquid helium. A typical stress-strain curve is shown in

Fig. 2. Because temperature rises occurred simultaneously with the

120
load drops, the specimens did not remain at 4 K. The strain required

to initiate serrated yielding decreased as the crosshead speed increased,

and the magnitude of the first detectable load drop also decreased with

crosshead speed, as shown in Fig. 3. No systematic effect of crosshead

speed on the yield strength was detected, Fig. 4.

The load drops were distributed in stress up to about 200 MPa,

which was over ten percent of the ultimate strength for this material.

A histogram showing the distribution of the magnitudes of the load drops

is shown in Fig. 5.

The maximum temperature rise observed was 52 K. A histogram of

the number of temperature rises in each of 11 temperature intervals,

measured using one thermocouple on one specimen strained to failure at

4 K, is shown in Fig. 6. Many load drops occurred with no indication

of a temperature rise at the thermocouple because in these cases the

temperature rises were localized in a small region of the specimen volume

away from the thermocouple. Assuming that the total number of temperature rises in a given temperature interval detected for all load drops

was proportional to the extent of specimen length attaining that temperature interval during a single load drop, a temperature profile of a

load drop was constructed, as shown in Fig. 7. The smooth curve is a

hand-drawn interpolation of the calculated stepped curve. The width at

half maximum of the calculated temperature profile was about 0.5 cm.

This agrees well with the value of specimen length involved in the load

drop obtained by visual observation of a specimen which was strained

through only one load drop. Fig. 3. By sprinkling this specimen with

iron powder, the region which was deformed during the load drop was

decorated because it contained more a' martensite than the part of the

129
specimen which had deformed uniformly. Visual observation of the

specimen surface before decoration also showed the severely deformed

region because its surface was visually deformed and the specimen was

necked. The deformed region appeared about 0.5 cm long, in agreement

with other observations.

The crosshead speed dependence of the strain at which the first

load drop occurred is consistent with expectations. The strain values

at which the first load drop occurred were different from those found by

Reed and Guntner [5]. This discrepancy is believed to be due to differences in chemical composition between the specimen materials of the

two studies, and confirms that the deformation process at 4 K in these

metastable austenitic stainless steel is quite sensitive to small

changes in chemical composition.

The widely accepted explanation of the phenomenon of discontinuous

yielding is adiabatic heating [1,2]. In this mechanism, the heat

produced in the specimen when it strains a small amount raises the

specimen temperature and lowers its strength enough for more deformation

to occur, and the process continues. The localized temperature instability at 4 K is caused by very low heat capacity and reduced thermal

conductivity. This cyclic process manifests itself in the stress strain

curve as a sudden load drop, accompanied by sudden specimen elongation.

The adiabatic heating model was tested by calculating the temperature

rise, the load drop, and the conditions for unstable plastic flow and

comparing the results with experimental data. The calculations follow

Basinski [1,2], but the discussion of unstable plastic flow differs in

detail. One assumes essentially adiabatic conditions, in which all

mechanical work is converted to heat, negligible conduction of heat away

130
from the deformed region, and negligible energy used in the generation

of lattice defects.

The temperature rise was calculated using the expression

T

/

max

CdT (1)

r.

1 m t

where P was the load on the specimen. Ax its length change, p its mass

density, v the specimen volume heated, C the specimen material specific

heat in -J/kg, T the temperature, the specimen temperature before

the load drop (= 4 K), and the maximum specimen temperature attained. Values for the parameters used in this and the following calculations are listed in Table 3. A graph of the specific heat over the

requisite temperature range was constructed using data from Ho, King,

and Pickett [6], Corsan and Mitchem [7] , and Jelinek and Collings [8].

The effect of the approximate nature of the specific heat curve used on

calculated temperatures is insignificant due to the strong temperature

dependence of the specific heat in the requisite temperature range.

From Eq. (1), the calculated value of T was 58 K, which agrees quite

max

well with the measured value, 56 K. The close agreement is probably

fortuitous due to the approximations involved. These include: assuming

that only a constant volume v is heated, and the temperature everywhere

else remains at 4 K, when actually a distribution of temperatures

through the specimen exists; and assuming that conditions are fully

adiabatic. The close agreement between measured and calculated T^^^

provides assurance that these two approximations are reasonably accurate.

131
The magnitude of the load drop was calculated from the expression

Aa = (8a/3T)AT + (9a/9c) Ac ( 2 )

where the first term on the right represents thermal softening and the

second accounts for strain hardening during the load drop, a denoting

stress and e strain. Note that Aa and 9a/9T are negative quantities,

while AT and Ac are positive, so that the thermal softening is opposed

by work hardening, as must be the case. If work hardening is ignored

(9a/9e = 0) we calculate Aa = -130 MPa for a temperature rise of 52 K.

Thus, thermal softening alone is more than enough to account for a

typical stress drop of 75 MPa. The appropriate value for the work

hardening coefficient 9a/9e is an average over strain and temperature.

If 9a/9a is approximated as the average over the strain range of the

load drops at 4 K, the calculated stress drop becomes 96 MPa, which is

closer to the typical value and within the observed range of stress

drops. The agreement between the calculated value (96 MPa) and the

observed value (typically 75 MPa) is considered adequate considering the

approximations involved.

The conditions for unstable plastic flow are given by

where k is the mechanical compliance of the load train, P is the load on

the specimen, and 9P/9T and 9P/9c are related, respectively, to 9a/9T

and 9a/9e by multiplying the latter by the specimen cross sectional

area. Table 3 gives a complete list of the meanings of the symbols.

This equation signifies that the load drop and heating continue up to

the temperature at which the specific heat C(T) is equal to

Pk/pv9P/9T • 1/(1 + k/£.9P/9c)> C(T) (3)

132
the heat produced for a small increment of strain divided by the mass of

heated specimen material. If work hardening is again ignored, we calculate a maximum specimen temperature of 58 K in good agreement with the

observed value 56 K. If work hardening is included using the approximation previously discussed for the work hardening coefficient, a maximum temperature of 44 K, which is considered adequate agreement with

experiment. It is concluded that comparison of the temperature rises

measured during load drops to those calculated using the magnitude of

the load drops, the load, the specific heat, the temperature dependence

of the yield strength, and the volume of the specimen participating in

the temperature rises confirmed adiabatic heating as the mechanism of

the load drops.

g The crosshead speed independence of martensite contents and yield

strength is consistent with the absence of deviations of specimen temperature from 4 K except during load drops. All these results indicate

that no dynamic effects were occurring, except during load drops.

The fact that no temperature rises were observed except during load

drops provides assurance that typical liquid helium temperature test

specimens and procedures do indeed yield 4 K data.

CONCLUSIONS

The following may be concluded from this study of the variability

of the yield strength and deformation behavior of A.I.S.I. 304L stainless

steel

.

1) Adiabatic heating is confirmed as the mechanism for load drops;

at higher strain rates, load drops occur at lower strains.

2) Small tensile specimens immersed in liquid helium remain at 4 K

during straining except during load drops.

3) During load drops, localized regions of the specimen deform.

133
REFERENCES

1. Z. S. Basinski, "The Instability of Plastic Flow of Metals at Very Low

Temperatures," Proc. Roy. Soc. A240 (1957) pp. 229-242.

2. Z. S. Basinski, "The Instability of Plastic Flow of Metals at Very Low

Temperatures," Austral. J. Phys. 1_3 (1960) pp. 354-358.

3. R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range 300 to

4 K," Adv. Cryo. Eng. 7 (1961) pp. 448-454.

4. R. P. Reed and R. P. Mikesell, "Low Temperature Properties of Selected

Copper Alloys", J. Mat. 2 (1967) pp. 370-392.

5. R. P. Reed and C. J. Guntner, "Stress-Induced Martensitic Transformations

in 18Cr-Ni Steel," Trans. Met. Soc. AIME 230 (1964), pp. 1713-1720.

6. J. C. Flo, G. B. King, and F. R. Fickett, "Low Temperature Specific Heat

of Two Stainless Steels," Cryogenics X8 (1978) pp. 296-298.

7. J. M. Corsan and N. I. Mitchem, "The Specific Heat of Stainless Steels

Between 4 K and 300 K," in Proceedings of the Sixth International Cryogenics Engineering Conference (IPC Science and Technology, Surrey,

England, 1976) pp. 342-344.

8. F. J. Jelinek and E. W. Collings, "Low-Temperature Thermal Expansion and

Specific Heat Properties of Structural Materials," in Materials Research

for Superconducting Machinery - Part 4 (1975). Available from NTIS,

Order Number ADA 019230.

9. D. T. Read, R. P. Reed, and R. E. Schramm, "Low Temperature Deformation

of Fe-18CR-8Ni Steel, in this volume.

134
LIST OF TABLES

Table 1

Table 2

Table 3

Chemical composition as supplied by the manufacturer of the AISI

304L alloy used in the present study.

Strain rates used, as calculated from crosshead speeds and effective gage length used in the present study, assuming no work hardening.

Parameters used in adiabatic heating calculations.

135
Table 1. Chemical computation, as supplied by the manufacturer,

of the AISI 304L alloy used in the present study.

Cr m Mo Ph Si

18.5 8.75 1.43 .29 .025 .31

S C N Cu

.013 .030 .061 .18 .14 bal

136
Table 2. Strain rates as calculated from crosshead

speeds and gage effective length used in the present

study, assuming no work hardening.

Crosshead Speed

(cm/mi n)

Strain Rate

(per second)

5 X 10'^ 2 X 10“^

2 X 10"^ 8 X 10"^

5 X 10'^ 2 X 10"^

137
Table 3. Parameters used in adiabatic heating calculations.

S.ymbol Meaning Value

P load 8 kN

AX specimen length change 1.27 X 10“

1 specimen length heated 5 X lO’^ m

A cross sectional area 1.2 X 10'^

V specimen volume heated = S, • A 6 X 10"® m

P specimen density 7.9 X 10"®

do/dJ thermal softening coefficient 2.5 MPa/K

9a/8e work hardening coefficient 1500 MPa

1C load train compliance 2.6 X 10"^

C (50K) specific heat at 50 K 88 J/kg K

m

m

kg/m^

m/N

138
LIST OF FIGURES

Figure 1

.

Figure 2.

Figure 3.

Figure 4.

Figure 5.

Figure 6.

Figure 7.

Figure 8,

Flat tensile specimen with thermocouple and superconductive

wire for detection of temperature rises.

Typical stress strain curve for AISI 304L at 4 K.

Strain required to initiate serrated yielding and magnitude of

the first detectable load drop, as a function of crosshead

speed

.

Yield strength at 4 K plotted against crosshead speed.

Histogram showing the frequency of stress drop magnitudes for one specimen strained to failure at 4 K,

Histogram showing the frequency of temperature rise magnitudes

recorded by thermocouple in one specimen strained to failure

at 4 K.

Profile of the temperature rise which occurs during a load drop,

as calculated from the distribution of temperature rises

occurring at one point. The smooth curve is a hand-drawn

interpol ation.

Photograph of a tensile specimen strained through one load

drop, decorated with iron powder.

139
Figure 1

.

Flat tensile specimen with thermocouple and superconductive

wire for detection of temperature rises.

140
STRESS

(MPa)

Figure 2. Typical stress-strain curve for AISI 304L at 4 K.

141
6

Figure 3

ok_ a

<oo

o' k. a.

CD

k_

5 -

3-

2 - 'ERROR BARS'

Indicate Relative

Magnitude of Load Drops x

I I I lO"* 10“^ 10“^ 10"' 10°

Crosshead Speed (cm/min)

Strain required to initiate serrated yielding and magnitude of

the first detectable load drop, as a function of crosshead

speed

.

14 .
Crosshead Speed (cm/min)

Figure 4. Yield strength at 4 K plotted against crosshead speed.

143

Yield

Strength

(KSi)
Magnitude of Stress Drop (MPa)

Figure 5. Histogram showing the frequency of stress drop magnitudes for one specimen strained to failure at 4 K.

144
28-

26-

24

= 22

TOI 20 E^ 18

a! 16

14-

2 12 -

I

A.I.S.I. 304L, 4K

Thermocouple Data

Total Temperature Rises Detected: 27

Total Load Drops: 101

6 -

4-

2 -

0 5 5 18 10 36 15 54 20 72 25 9 31 08 3626 41 44 46 62

5 18 10 36 15 54 20 72 25 9 31 08 36 26 41 44 46 62 51 8 ATTK)

Figure 6. Histogram showing the frequency

recorded by thermocouple in one

at 4 K.

of temperature rise magnitudes

specimen strained to failure
Temperature

Rise

(K)

Figure 7. Profile of the temperature rise which occurs during a load drop,

as calculated from the distribution of temperature rises

occurring at one point. The smooth curve is a hand-drawn

interpolation.

14G
Figure 8. Photograph of a tensile specimen strained through one load

drop, decorated with iron powder.

147

LOW TEMPERATURE DEFORMATION

OF Fe-18Cr-8Ni STEEL

National Bureau of Standards

149

LOW TEMPERATURE DEFORMATION OF Fe-18Cr-8Mi STEEL*

D. T. Read, R. P. Reed, and R. E. Schramm

Fracture and Deformation Division

National Bureau of Standards

Boulder, CO 80303

*Work supported by DOE/OFE. Contribution of NBS, not subject to copyright.

151
ABSTRACT

The flow strengths of an Fe-18Cr-8Ni austenitic alloy were measured

over the temperature range 4-300 K. This alloy did not transform

martensi tical ly on cooling but did transform to two martensitic products

a' (body-centered-cubic, ferromagnetic) and e (hexagonal -close-packed,

non-ferromagnetic) during deformation. Volume percentages of the

austenite and martensite phases which formed during deformation and

stacking fault and twin fault probabilities of the austenite were measured as a function of strain at 4 K using x-ray and magnetic techniques. The flow strengths were measured over the temperature range 4-

320 K at constant strain rate. The temperature dependence of the yield

strength was anomalous, with a peak near 200 K and a minimum near 100 K,

followed by steadily increasing strength to 4 K. These features have

been attributed to the formation during deformation of the martensitic

phases. Similar to previous results, the hexagonal -close packed phase

(e) increased to a maximum of about 20 percent at about 10 percent

strain, then decreased with additional deformation. The body-centered

phase (a‘) increased continuously as a function of strain. Stacking

fault probability decreased as the e martensite content increased.

Martensite contents and fault probabilities were found to be independent

of strain rate at 4 K. Discussion is presented on the role of the

martensitic transformation in affecting deformation of austenitic alloys.

152
INTRODUCTION

Despite widespread use and study of metastable austenitic stainless

steels at cryogenic temperatures , the tensile behavior of these alloys

under various conditions of temperature, strain rate and cold work are

not understood. These steels are extremely complex. Lower nickel, FeCr-Ni austenitic steels have a magnetic transformation from paramagnetic

to antiferromagnetic states with the Ndel temperature usually near 60 K.

Higher nickel, Fe-Cr-Ni austenitic steels have a more complex magnetic

transformation; the low temperature state is now thought to represent a

spin-glass type condition of localized ferromagnetic spins, dispersed

randomly. Most AISI 300 series grades of austenitic stainless steels do

not transform martensi tical ly on cooling (i.e. temperature < 0 K) but

do transform on deformation at low temperatures to two martensitic

products: (1) a' (body-centered cubic, ferromagnetic) and (2) e (hexagonal close packed, either paramagnetic, antiferromagnetic or spinglass magnetic state). Furthermore, in the Fe-18Cr-8Ni grades the

stacking fault energy (y) is low at room temperature 20 erg/cm ) and

decreases toward zero at low temperatures. This low stacking fault

energy insures the formation of many stacking faults during low temperature deformation.

Some studies have reported anomalous flow strength behavior for FeCr-Ni base austenitic steels at low temperatures, Breedis and Robertson^

found that as the (a') temperature (241 K) of Fe-16Cr-12Ni single

crystals was approached from higher temperatures the critical resolved

shear stress decreased to zero. Deviations from the normal temperature

dependence were apparent up to about 100 K above the M^ temperature.

Stone and Thomas also reported decreased critical resolved shear stress

153
at low temperatures (180-190 K) for Fe-15Cr-15Ni single crystals. Sato

3

et al conducted a series of tensile tests on Fe-18Cr-14Ni single crystals in the temperature range 170-430 K. Their tensile measurements

indicate a clearly defined temperature (243 K) at which the flow stress

begins to decrease. This decrease was attributed to hexagonal martensite formation, and 243 K was described as the •^ij(e) temperature. From

examination of specimen surface morphology, Sato et al concluded that

below 243 K the dominant shear system was {111} <112> and at higher

temperatures the more conventional {111} <1T0> system was observed.

4 Higo et al , working with perhaps the same Fe-18Cr-14Ni single crystals,

had earlier found the same flow-strength temperature trends and postulated {111} <112> shear systems in the austenite.

5 Measurements by Ilichev and Skibina, described by Verkin , from 4

to 300 K on Fe-18Cr-8Ni, Fe-18Cr-10Ni and Fe-18Cr-20Ni single crystals

indicate the possibility for three separate dips of the yield-strength

temperature dependence.

Measurements indicating strength reductions at lower temperatures

for polycrystal 1 ine Fe-Cr-Ni austenitic stainless steels have been

performed by Mirzayev et al^ and Suzuki et al^’^. Tensile yield strengths

of the alloys Fe-18Cr-10Ni and Fe-18Cr-13Ni alloys were measured from 77

to 670 K by Mirzayev et al . Above the M^(a') temperature, a reduction

of yield strength of about 20 percent was observed from higher temperature

values. Contrasted to the single crystal results, as the M^(a') temperature was approached, the yield strength again started to rise. The

resultant yield strength temperature dependence was a dip (of about 20

percent) between the M^(a') temperature and about 50 to 100 K above this

temperature. This decrease of flow strength at low temperatures was

154
7 8 attributed to the transformation to e martensite. Suzuki et al ’ also

have conducted tensile measurements between 77 and 400 K on polycrystalline Fe-18Cr-9Ni alloys. Within the temperature interval where the flow

stress is decreasing with decreasing temperature they have identified

two distinct flow stresses. The lower stress corresponds to the proportional limit and to the stress at which e martensite forms (although

the technique for e martensite observation is not clear). The higher

stress corresponds to the stress at which a' martensite is first detected magnetically. The yield strength (0.2% offset) corresponds to

the temperature dependence of the lower, e-associated stress.

On the other hand, most research on TRIP steels, in which flow

stress decreases at decreasing temperatures have also been identified,

9 correlate flow strength with a' martensite content. Olsen and Azrin ,

working with a polycrystalline Fe-9Cr-8Ni -2Si -0. 2 , 0.3C alloy have

associated the flow strength (0.2%) with the stress necessary to form

0.01 volume fraction a'. Below the temperature 320 K) at which the

strength begins to decrease, the temperature dependence of both stresses

coincide.

This paper measures the low temperature dependence of the flow

strengths for a standard, commercially available AISI 304L. In this

alloy the flow strengths were also found to dip at intermediate low

temperatures . Associated magnetic and x-ray data are presented to

identify the controlling transformation for low temperature deformation

of this alloy.

MATERIAL

The material was a 2.5 cm (1 inch) plate of AISI 304L stainless

steel obtained in the mi 1

1 -anneal ed condition. Its chemistry is shown

155
in Table 1. Flat tensile specimens in the longitudinal orientation with

a gage section 2.5 mm (0.1 inch) thick by 4.8 mm (.19 inch) wide by 30

mm (1.2 inch) were used for all tests. The nominal grain diameter was

45 ym and the hardness Rockwell B-79.

After machining, tensile specimens were annealed in vacuum, at

1000°C for 1 hour to remove possible cold work. The specimens were

removed from the furnace still under vacuum and allowed to cool. These

specimens were used for measurements of yield strength versus temperature. After testing to one percent strain, the specimens were reannealed under the same conditions and retested. Some specimens were

tested and annealed a total of four times. Grain sizes ranged from 40

to 60 ym after the initial anneal to 60 to 80 ym after four anneals.

Some flow strength (0.2% offset) dependence on grain size was detected

at temperatures lower than 100 K, but little dependence was apparent

above 100 K. All data points, independent of grain size (40-80 ym),

are included in this report.

The observed M^(a') and M^(e) temperatures were below 4 K for this

alloy. A calculated transition temperature of M^(a') = 53 K was obtained, following Larbalestier and King^^ who used the work of Eichelman

and Hull and Hammond on austenitic (y->a') stability equations for Fe-CrNi steels. The calculated Mj(a') temperatures, using the formula developed by Angel for 30 to 50 percent tensile deformation, is 288 K.

TECHNIQUES AND APPARATUS

The tensile tests were carried out using a standard tensile machine

with the capability for variable crosshead speeds. A cryostat arrange1

2

ment developed by Reed was used for tensile tests at 4 K. Displacement
1

3

gages as described by Reed and Mikesell were used to measure specimen

strain. This system was modified for use from 4-320 K with the addition

of a reservoir for liquid nitrogen or liquid helium on the upper pull

rod, heaters on both the upper and lower specimen grips, and electrical

resistance thermometers (platinum for temperatures above 76 K, carbon

for lower temperatures) . The apparatus is shown in Fig. 1. Liquid

level sensors, variable-output power supplies for the heaters, and

amplification and readout electronics for the thermometers were used

with the equipment shown in Fig. 1. All x-ray and magnetic measurements

were conducted at room temperature. The body-centered cubic, ferromagnetic, a' martensite, was measured using both x-ray diffraction

and magnetic techniques. The hexagonal close packed, e martensite was

only measured by x-ray techniques.

Magnetic measurements were carried out using a commerciallysupplied apparatus incorporating a permanent bar magnet on an adjustable

helical spring. From these measurements the a' concentration can be

calculated, using a previously established relationship between a' and

magnetic susceptibility .

The x-ray measurements were conducted as follows. The integrated

intensity of x-ray diffraction lines is an indicator of the volume

fraction of each crystal lographic phase present. Assumming random

. 15-17

orientation :

[A FxB (

1 )

where

stant, m^_

xp

= diffracted intensity for line x of the 3-phase,

= multiplicity factor, LP = Lorentz-polarization

A = confactor =

157
2 2

1 + cos 29/sin 0-cos 9 where 9 = Bragg angle for line x, = Xp

-2M

structure factor for line x of phase 3, e = temperature factor, v„ =

volume of unit cell of phase 3 and = volume of phase 3.

3

For convenience the above equation can be abbreviated

I x3 ( 2 )

where is a constant for each diffraction line of phase 3-

The diffraction lines used for the three phases are: (200)y>

(200)a‘ and (lOTl)c. The radiation was Co Ka. Hence, the calculated

values of K and K relative to K , are

y £ a

K = 4.5K , (3a)

£ a and

K = 2.3K , (3b)

Y a

where y refers to the austenite phase. Lattice parameter values used to

o o o obtain these ratios were 2.866 A, for a'3.589 A for a and a = 2.54 a

° 1

5

and c = 4.15 A for e . Dispersion by the K electrons caused corrections

to the atomic scattering factor (contained within F^^) estimated as

1

8

-3.9% for all phases . Estimates of temperature factors and atomic

1 fi scattering factors were from Cullity

With the volume of a' known from magnetic measurements, the volume

of e can be obtained from

''c " " '(lOilJ /4.5 K^, = [I (lOTl )y I (200)^,V^,/4. 5] (4)

158
and the volume of retained austenite is easily obtained using

Discrete intensity measurements were made at 0.04° 20 intervals on

a diffractometer across entire diffraction lines. The background was

computed from a straight line fitted to average intensity values on

either extreme of the lines. Numerical integration by the trapezoidal

rule provided the diffraction peak area. Three separate independent

background and intensity measurements were averaged.

Errors in the magnetic measurement of a' volume are about + 2%

and the x-ray measurements are estimated to have a precision of about

+ 6 %. Analysis of repeated measurements of e at two strains led to a

probable error in y less than + 7%.

Stacking faults on the (111) plane in austenite result in a decreased separation of the peak positions of the (111) and (200) diffraction lines. A quantitative measure of the probability of these defects,

a, comes from comparing this line separation between annealed and coldworked specimens^

(20200 ^^lll^CW ^^®200 ^®111^AM

" ^^^"^^200 tan0^^.j)a

Peak positions were determined by measuring diffracted intensity at five

to seven 20 positions centered around the peak, fitting this data to a

parabola and calculating its maximum. Probable imprecision on stacking

fault probability is + 4 x 10~^.
Twin faults produce an asymmetry of the diffraction line as measured by the difference between the center of gravity and the position

20

of the peak maximum . Here

ACG(°20)^^^ - ACG(°20)2qq = (11 tan0^^^ + 14.6 tan02QQ)3 (7)

where ACG = center of gravity - peak maximum and 3 = twin fault probability. Because of instrumental asymmetry, the values of ACG need

21 correction by comparison with an annealed standard . The center of

22

gravity calculation came from the profile measurements noted above . _3

An error estimate of 3 is + 7 x 10 .

RESULTS AND DISCUSSION

Temperature Dependence of the Flow Strength

The temperature dependence of the flow strength at 0.2 percent

strain (corresponding to the conventional yield strength) is shown in

Fig. 2. There is a pronounced decrease in flow strength beginning on

cooling at about 200 K. A minimum is reached near 100 K, about 15

percent lower than the 200 K flow strength. Below 100 K, the flow

strength at 0.2 percent strain rises very steeply to 4 K. Although

there is some data scatter, the possibility for the existence of three

dips in the temperature dependence of the flow strength reported by

5 Verkin seems remote.

The anomalous flow strength behavior at low temperatures is primarily confined to low plastic strains. The temperature dependence of

the flow strength at 1.0 percent strain is presented in Fig. 3. While

ICO
the trend continues with the temperature dependence changing near 200 K

on cooling, there is no longer a decrease in strength from 200 to 100 K.

At higher strains the plateau of strength versus temperature observed

between 100 to 200 K becomes more obscure.

As discussed earlier, other studies have identified either the

5-8

hexagonal close packed (c) phase or the body-centered cubic (ct')

9

phase as the transformation product controlling the early stages of

deformation in polycrystall ine austenitic alloys. Any model must explain two phenomena: (1) the abrupt decrease of the flow strength at

about 200 K with decreasing temperature and (2) the abrupt increase of

the flow strength near 100 K with decreasing temperature. X-ray and

magnetometry experiments are now in progress to measure the amounts of

austenite and c and a' during the early stages of deformation at 200,

100 and 4 K. These measurements are not easy since the amounts of e and

a' are of the order of 1 percent.

Dependence of Martensitic Transformations on Plastic Deformation at 4 K

The volume percentages of austenite, e and a' as a function of

plastic deformation at 4 K were estimated using x-ray (c) and magnetic

measurements (a'). The results are presented in Fig. 4.

There is a linear dependence of percent a' formed at low temperatures as a function of strain. If the data of earlier measure9 15 23-24

ments *

* on austenitic Fe-Cr-Ni alloys at 4 and 77 K are evaluated, the same conclusion is reached. However, there is marked quantitative discrepancy in the dependence of a' formation on deformation that

cannot be adequately explained through use of current stability equations for or temperatures. There is no apparent correlation

strain between either calculated M^(a') or M^(a') temperatures and

percent a' formed at constant strain.

161
The hexagonal close-packed martensite phase initially increases

with increased strain, then decreases at higher strains. The decrease

of e has led many to conclude that the e phase transforms to a'. Presumably at higher strains (above approximately 10 percent elongation)

the effects of higher flow stresses and austenite partitioning combine

to retard the austenite to e transformation.

The dependence of the stacking fault and twin fault probabilities

on plastic deformation at 4 K, shown in Fig, 5, have not been previously reported. The stacking fault probability decreases at low

strains, reaches a minimum at approximately 10 percent strain and increases at higher strains. These data correspond very well with the

dependence of the volume percent hexagonal phase on strain; the presence

of stacking faults and hexagonal phase are inversely related. A net

reduction in the number of stacking faults present in a deformed material implies that these existing stacking' faults are forming hexagonal

martensite (or, less likely, a' martensite) faster than they are being

generated. However, at higher strains, the process is reversed; presumably the newly formed stacking faults are not acting as nuclei for

either e or a' martensite and, thus, increase in density. Furthermore,

the rate of e formation is now exceeded by the rate of formation of a'

martensite from e.

Deformation twin fault density is shown in Fig. 5 to remain constant with strain. The magnitude of twin fault density scatters about

3 the value for the annealed condition (3 x 10 ), indicating the absence

of measurable formation of deformation twins in this alloy. However,

examples of twin fault formation have been reported in an Fe-18Cr-7Ni0.18C steel deformed 20 percent at 380

1G2
CONCLUSIONS

The following may be concluded from this study of the temperature

dependence of flow strength and deformation behavior of an Fe-18Cr-8Ni

stainless steel:

(1) The temperature dependence of the flow strength is anomalous,

with a peak near 200 K and a minimum near 100 K.

(2) The volume percent of strain-induced body-centered cubic

martensite is linearly dependent on strain and independent of strain

rate at 4 K.

(3) The volume percent of strain-induced hexagonal close-packed

martensite at 4 K reaches a maximum near 10 percent strain.

(4) The number of strain-induced stacking faults at 4 K reach a

minimum near 10 percent strain.

ACKNOWLEDGEMENTS

The authors wish to express their thanks to David Burkhalter, who

conducted many of the tensile tests in this study. The work was supported by the Office of Fusion Energy, Department of Energy.

163
REFERENCES

1. J. F. Breedis and W. D. Robertson, "Martensitic transformation and plastic deformation in iron alloy single crystals," Acta Met. 1J_ (1963), pp.

547-559.

2. G. Stone and G. Thomas, "Deformation induced alpha and epsilon martensites in Fe-Ni-Cr single crystals," Metall. Trans. ^ (1974), pp. 2095-

2102.

3. A. Sato, Y. Sunaga and T. Mori, "Contribution of the y-j-e transformation

to the plastic deformation of stainless steel single crystals," Acta Met.

^

(1977), pp. 627-634.

4. Y. Higo, F. LeCroisey and T. Mori, "Relation between applied stress and

orientation relationship of a' martensite in stainless steel single

crystals," Acta Met.

^

(1974), pp. 313-323.

5. B. I. Verkin, "The development of cryogenic materials science in the

USSR," Advances in Cryogenic Engineering, Vol . 24, R. P. Reed, A. F. Clark, K. Timmerhaus, Editors, Plenum Press, New York (1977).

6. D. A. Mirzayev, Yu. N. Goykdenberg, M. M. Shteynberg and S. V. Rushchin,

"The elasto-plastic strain effects in alloys with low stacking fault

energy at temperatures above M ," Phys. Metals Metal loy. No. 6 (1973)

pp, 81-87.

7. T. Suzuki, H. Kojima, K. Suzuki, T. Hashimato, S. Koike and M. Ichihara,

"Plastic deformation and martensitic transformation in an iron-base

alloy," Sripta Met. 1_0 (1976), pp. 353-358.

8. T. Suzuki, H. Kojima, K. Suzuki, H. Hashimato and M. Ichihara, "An experimental study of the martensite nucleation and growth in 18/8 stainless

steel," Acta Met.

^

(1977), pp. 1151-1162.

9. G. B. Olson and M. Azrin, "Transformation behavior of TRIP steels,"

Metall. Trans. A. 9A (1978), pp. 713-721.

10. D. C. Larbalestier and H. W. King, "Prediction of the low- temperature

stability of type 304 stainless steel from a room temperature deformation

test," Proc. Intn. Cryogenic Engr. Conf . , IV, Eniehoven, Netherlands (May

1972).

11. T. Angel, "Formation of martensite in austenitic stainless steels," J.

Iron Steel Inst. UT. (1954), pp. 165-174.

12. R. P. Reed, "A cryostat for tensile tests in the temperature range 300 to

4 K," Adv. Cryo. Eng. 7 (1961), pp. 448-454.

13. R. P. Reed and R. P. Mikesell, "Low temperature properties of selected

copper alloys," J. Mat. 2 (1967), pp. 370-392.

1G4
14. R. P. Reed and R. P. Mikesell, "The stability of austenitic stainless

steels at low temperatures as determined by magnetic measurements," Adv.

Cryo. Eng. 4^ (1960), pp. 84-100.

15. R. P. Reed and C. J. Guntner, "Stress-inducted martensitic transformations in 18Cr-8Ni steel," Trans. Met. Soc. AIME 230 (1964), pp. 1713-

1720.

16. B. D. Cullity, Elements of X-ray Diffraction , Addison-Wesley Pub. Co.,

Inc., Reading, MA (1956), pp. 389, 395, 460, 475 and 477.

17. B. L. Averbach, M. F. Camerford and M. B. Bever, "An x-ray method for the

determination of beta phase in a titanium alloy," Trans. Met. Soc. AIME

215 (1959), pp. 682-685.

18. R. W. James, The Optical Principles of the Diffraction of X-rays , G.

Bell & Sons, Ltd., London (1954), p. 608.

19. R. E. Schramm and R. P. Reed, "Stacking fault energies of fee Fe-Ni

alloys by x-ray diffraction line profile analysis," Met. Trans. 7A (1976),

pp. 359-363.

20. J. B. Cohen and C. N. J. Wagner, "Determination of twin fault probabilities from the diffraction patterns of fcc-metals and alloys," J. Appl

.

Phys. ^(1962), pp. 2073-2077.

21. E. N. Aqua and C. N. J. Wagner, "An x-ray diffraction study of the effect

of rhenium on the occurence of faulting in refractory alloys," Office of

Naval Research, Tech Rep. No. 11, Contract N0N0R609(43) , 1964.

22. R. E. Schramm, "Corrections and calculations on an x-ray diffraction line

profile: a computer program," National Bureau of Standards, Boulder, CO.

23. G. Blanc, R. Tricot and R. Castro, "Transformations martensi tiques dans

les aciers inoxydables austeni tiques Fe-Cr-Ni," Mem. Sci . Rev. Metall.

70 (1973), pp. 527-541

.

24. H. C. Fielder, B. L. Averbach and M. Cohen, "The effect of deformation on the martensitic transformation in austenitic stainless steel," Trans. ASM

47 (1954), pp. 267-286.

25. L. Remay and A. Pineau, "Observation of stacked layers of twins and e martensite in a deformed austenitic stainless steel," Metall. Trans. ^

(1974), pp. 963-965.

165
Table 1. Chemical composition, as supplied by

facturer, of the AISI 304L.

the manu

Cr Ni Mn Mo Ph Si

18.5 8.75 1.43 .29 .025 .31

S C N Co Cu Fe

.013 .030 .061 .18 .14 bal

16G
LIST OF FIGURES

Fig. 1 Apparatus used for variable temperature tensile tests.

Fig. 2 Yield strength (0.2 percent strain) of AISI 304L in the temperature

range 4-320 K. The minimum near 100 K and the peak near 200 K are anomalous features.

Fig. 3 Flow strength at 1 percent strain of AISI 304L in the temperature

range 4-320 K.

Fig. 4 Martensite and parent phase contents as functions of strain at 4 K for three crosshead speeds, with hand-drawn trend lines.

Fig. 5 Stacking and twin fault probabilities as functions of strain at 4 K for three crosshead speeds, with hand-drawn trend lines.

167
Fill Tube

Displacement Gage

Specimen

Liquid Nitrogen

or Liquid Helium

Heaters

Thermometers

VARIABLE TEMPERATURE TENSILE TEST APPARATUS

Fig. 1 Apparatus used for variable temperature tensile tests.

IGo
400

O)c

O)

250

200

AISI 304L

Laboratory-Annealed

_J_

100

_l I •

200 300

Temperature (K)

CSI d

Fig. 2 Yield strength (0.2 percent strain) of AISI 304L in the temperature

range 4-320 K. The minimum near 100 K and the peak near 200 K are anomalous features.

169
600

-80

-70

o

SPo

00S

400 -

-60

300 -

A.I.S.I. 304L

Laboratory-Annealed

-40

200

100 ±

200

Temperature (K)

- 30 300

Fig. 3 Fluw strength at 1 percent strain of AISI 304L in the temperature

range 4-320 K.

170

Flow

Strength,

1.0 %

Offset

(ksi)
Fig. 4 Martensite and parent phase contents as functions of strain at ^ K for three crosshead speeds, with hand-drawn trend lines.

171
TWIN

FAULT

PROBABILITY

(xlO-^)

STACKING

FAULT

PROBABILITY

(xlO-")

Fig. S Stacking and twin fault probabilities as functions of strain at 4 K fur three crosshead speeds, with hand-drawn trend lines.

172
Physical Properties of Stainless Steels

173
(
ELASTIC-CONSTANT VARIABILITY IN STAINLESS STEEL 304

National Bureau of Standards

Elastic-Constant Variability in Stainless Steel 304

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

M. V. Frederick

Electromagnetic Technology Division

National Bureau of Standards

Boulder, Colorado 80303

M.. W. Austin

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

Variability of elastic constants in stainless steel 304 was determined by measuring longitudinal and transverse ultrasonic velocities

in twenty samples acquired randomly. Three kinds of variations - -

sample-to-sample, directional within a sample, and repeated measurements

on a single sample - - are reported for four elastic constants: the

bulk modulus. Young's modulus, shear modulus, and Poisson's ratio.

Because of surprisingly small variations, 1 percent or less, the

principal problem became measurement sensitivity and reproducibi 1 i ty.

To overcome this problem, an inexpensive high-resolution measurement

system was devised using general -purpose equipment augmented with a very

simple impedance-transformi ng amplifier and a broadband detector. With

this system the often-reported , troublesome transit-time correction disappeared. Effects due to frequency and directionality were negligible.

*Contri bution of the U. S. Government, not subject to copyright.

177
I. INTRODUCTION

For stainless steel 304, the presently reported study sought to

determine the variability of elastic constants such as the bulk modulus.

Young's modulus, shear modulus, and Poisson's ratio from sample to

sample, directionally within a sample, for repeated measurements, and as

a function of sound-wave frequency.

Elastic constants of technological materials fill many needs. They

become vital in problems such as load-deflection, thermoelastic stress,

elastic instability, and fracture mechanics. Empirical correlations

occur between elastic constants and technologically important properties

such as strength, hardness, and wear. Because they relate intimately to

both the interatomic potential and the lattice-vibrational spectrum,

elastic constants relate in turn to various fundamental solid-state

phenomena such as diffusion, theoretical strength, and lattice defects.

Since elastic constants occur as parameters in thermodynamic equations

of state, well-defined relationships exist between elastic constants and

free energy, specific heat, and thermal expansivity. Knowledge of the

elastic constants allows one to compute the elastic Debye temperature,

which equals the specific-heat Debye temperature at zero temperature.

The Debye temperature characterizes crystalline solids better perhaps

than any other single parameter.

Although usually expected, random variability of physical and,

especially, mechanical properties of technological materials is seldom

character! zed quantitatively. A recent study^ considered the mechanicalproperty variability among fifteen product forms from a single lot of

stainless steel 304. Variability arises from three principal causes:

chemical composition, thermomechanical treatment, and testing methods.

170
The present study used a single testing method, thus obviating the third

cause.

Five types of stati stical -data samples suggest themselves for a

property variability study: (1) lot-to-lot; (2) sample-to-sample, within

lot; (3) within sample, directional; (4) within sample, repeated measurements; (5) within sample, various test methods. The present study

considers statistical samples (1), (3), and (4), and the results of the

study show that studying statistical sample (2) seems seldom worthwhile

in the case of elastic properties.

The most accurate elastic constants, C, usually follow from determining ultrasonic velocities, v, and using the general relationship

2

C = pv , where p denotes mass density. The present study used a simple

pulse-echo method to determine velocity v = 2£/t, where I denotes the

specimen length and t denotes the round-trip time for an ultrasonic

pulse between two fl at-and-paral lei specimen faces. As described below,

high accuracy in the present measurement results from four features:

narrow pulse width, broad bandwidth, unrectified echoes, and cycle-tocyle time measurements using consecutive echoes.

Ultrasonic-velocity determinations of elastic constants in stainless2 steel -304-type alloys have been reported previously. In 1960, Fahey

reported ultrasonic elastic constants for stainless steel 303 at four

3

temperatures between ambient and liquid-helium. Haskins reported a

similar study for stainless steels 303 and 321 down to 77 K and 23 K,

respectively. Armstrong and Eash^ used a thin-rod resonance method to

measure ultrasonic velocities in stainless steel 303 between 4 K and 320

K. Ledbetter, Naimon, and Weston reported the 4 to 300 K ultrasonic

velocities for four stainless steels: A286, 304, 310, and 316. Hammond,

179
r Moyer, and Brinkman reported ultrasonic elastic constants for stainless

304 between 300 and 1000 K. Bates^ recently correlated ultrasonic shear

modulus with irradiation-induced swelling in stainless steel 316, and

he reported the effect of Mo, P, and Si compositional variations on the

elastic constants.

The present report represents one of the first detailed studies of

g elastic-constant variability for a technological material. Richards

reported a variability study of several static and dynamic experimental

methods for determining elastic moduli in beryl 1 ium-copper rod.

II. EXPERIMENT

A. Materials

Twenty samples of stainless steel 304 were obtained effectively

randomly from various sources in various forms, including bar, pipe, and

plate. No laboratory treatments were performed on the alloys; they were

tested in their as-acquired conditions. Table I gives chemical composition,

hardness, and mass density.

U1 trasonic-velocity specimens were prepared by grinding flat and

parallel surfaces (within 3 pm). Specimens studied for directionality

variations were ground as rectangular prisms (Fig. 1).

B. Electronic system

The studied materials exhibit some dispersion, as shown by the

pulse-echo pattern in Fig. 2. Since the specimen heavily loads the

quartz transducer, the bandwidth of repeated echoes is limited mainly by

the impedance-matching network connected to the transducer. Consistent

with the common knowledge that time resolution varies inversely with the

bandwidth of the measuring system, a wide-bandwidth system was devised.

The bandwidth was about 20 times that of commercially available measurement systems used previously in our laboratory. Figure 3 shows the

180
circuit diagram of this new system. Note the absence of a variable

tuning circuit. Note also that the quartz crystal is loaded by a very

low capacitance cable and an FET source follower. This arrangement

allows a good match to the high output impedance of the quartz transducer capacitance without sacrificing bandwidth. This approach to

receiver design allowed the use of a low-power general -purpose oscillator as a signal source.

C. Velocity measurements

Velocities were determined by measuring the time difference between

consecutive echoes in a pulse-echo pattern using a dual -channel sweepdelayed time-interval -microprocessor-equipped oscilloscope with typically

a 0.001 ys resolution. Thus, for a 5 ys transit time the best resolution

in the velocity is 2 parts in 10,000. Neglecting density error this

amounts typically to a 0.04 percent error in the elastic constant. This

resolution can be improved, but it serves adequately in most cases. It

becomes more efficient to 'clean up' the echoes so they can be interpreted

and measured more accurately. And this was achieved with the system

described above. Figure 2 shows an oscilloscope display of the pulse

and the first two echoes. Time measurements were made between the

crests of leading cycles in consecutive echoes. Since the leading cycle

attenuates strongly, care was needed when measuring between nonconsecutive echoes to guarantee a correct cycle matchup.

D. Transit-time error

9-1

1

Many authors have reported a transit-time error due to the

transducer and bond thicknesses. Previous publications from our laboratory corrected for this error by either a two- transducer method or by

a multiple-specimen-length method.
The new measurement system eliminates the transit-time error,

identifying it as an artifact due to attenuation of an echo's leading

cycle and accumulation of spurious trailing cycles due either to the

specimen and the transducer or to the electronics.

Attempts to make a transi t-time-error correction using the procedure

suggested by Neighbours, Bratten, and Smith^^ using four differentlength specimens always resulted in corrections corresponding closely to

an integral number of cycles. This was confirmed at quartz-crystal

transducer fundamental resonance frequencies of 4, 5, 7, and 10 MHz for

the longitudinal case and at 4 and 8 MHz for the transverse case.

Correction attempts using a two-transducer method gave less welldefined echoes. But even these could be measured reasonably well by

cycle-cycle overlap to give zero transit-time error.

Perhaps the most convincing evidence for zero transit-time error is

shown in Table II, which shows velocities determined on a single specimen

for a series of 180 measurements. For each measurement the transducer

seal was broken and remade.

The careful reader has probably already noted that the conclusions

given here concerning transit-time error are essentially a rediscovery

1

2

of those given 20 years ago by Eros and Reitz .

III. RESULTS

Table II shows the effects of repeated measurements. Each entry in

Table II represents the arithmetic average of 10 measurements, each with

a different bond between specimen and transducer. Table II also shows

effects of using transducers with different fundamental frequencies: 3,

5, and 7 MHz.

182
Table III shows effects of changing frequency in another way. A

single 5 MHz transducer was driven off its resonant frequency between

2.5 and 7 MHz.

Table IV shows for a typical case the variation of velocity with

direction within a single specimen, from a 3.8-cm 304L plate. Here, 001

denotes the rolling direction and 100 denotes the short-transverse

di recti on

.

Table V gives the main results of the study--the longitudinal and

transverse sound velocities together with five derived elastic constants : ^ ^

longitudinal modulus =

(

1 )

shear modulus = G

=

bulk modulus = B

=

Young's modulus = E

=

and Poisson's ratio = v

=

CM

+-> a. > (2)

P (3)

3 py2 - I

’ (4)

- ^4^(4 - ’ (5)

together with their standard deviations.

IV. DISCUSSION

Results in Table II show that longitudinal and transverse ultrasonic

velocities in commercial stainless steels can be measured to an imprecision

of about 1 part in 1000, or 0.1 percent. Using familiar error-propagation

formulas, this translates approximately to imprecisions of 0.2 percent

in C^, 0.2 percent in G, 0.3 percent in B, 0.5 percent in E, and 0.5
percent in v. While such levy imprecisions are claimed often in singlecrystal experiments, the authors are unaware of a similar claim for

commercial -a Hoy polycrystals.

Measuring between different pairs of echoes (1 and 2, 2 and 3, 3

and 4 in Table II) does not affect the measured velocity beyond the

present resolution capability, typically 0.05 percent for v^ and 0.03

percent fo v^. This means the first echo is unaffected by the proximity

of a large pulse. (Wider pulse widths may negate this conclusion.)

Changing measurement frequency between 3 and 7 MHz does not affect

the measured velocity beyond the resolution capability, as shown in

Table II. This means that dispersion effects are either absent or quite

small, and except for familar adiabatic-isothermal corrections, elastic

constants measured at high frequencies apply also to static cases.

The 180 measurements summarized in Table II represent 180 different

transducer-specimen bonds as well as 6 different transducers. Thus, the

effects of changing transducers and bonds are also negligible. These

results show conclusively that transit-time errors discussed above

(section II. D) are nonexistent.

The only 'tuning' in the present measurement system consists of

changing the frequency. This was sometimes useful for improving the

echo-pattern quality, best-quality echo patterns being obtained at

frequencies between 2.5 and 7.5 MHz. For v^ the effect is nonmeasurable; v^ changes smoothly only about 0.2 percent over this frequency

range. Thus, the present measurement system responds reproducibly over

a moderately large frequency range. This contrasts sharply with our

laboratory's previous experience with measurement systems that contained

tuning (impedance-matching) circuits that gave measurable echo patterns

only over a narrow frequency range, about 0.5 MHz or less.
Directional variations of elastic constants v/ithin a sample are

quite small, as shown in Table IV. Other specimens studied similarly

also showed small directional variations. This means that textures are

nearly nonexistent in these alloys because single-crystal elastic constants for stainless steels^"^ show a Zener anisotropy 2C^^/(C-|-j - C-j^)

between 3 and 4, which would cause measurable elastic-constant variations if crystallites were oriented nonrandomly. This lack of texture

may be ascribed to the thermomechanical treatment: annealing after hot

deformation. Textures caused by cold deformation are less easily removed by annealing.

The most important, and perhaps most surprising, results of the

present study are shown in Table V, which gives the lot-to-lot variability of the sound velociites and the elastic constants. Velocity

variations are quite small, 0.3 percent for v^^ and 0.6 percent for v^.

Coupled with a 0.6 percent density variations, this leads to elasticconstant variations of 0.9 percent for E, 1.2 percent for v, and 1.4

percent for both B and v. All these are niuch smaller than variations

implied by elastic constants reported in the literature for these alloys.

Thus, reported variabilites in these elastic constants must arise from

experimental measurement errors, not from the material itself. Considering differences among these alloys that must exist in composition,

microstructure, grain size, and carbon distribution, the low elasticconstant variability is indeed surprising. And it would be interesting

to study this problem in other types of alloys: aluminum, copper, ferritic steel, and others.

V. CONCLUSIONS

The following principal results and conclusions emerge from the

present study:

185
1. A new pulse-echo elastic-constant measurement system permits

both longitudinal and shear velocity measurements to be made within

0.1 percent on materials such as commercial stainless steels.

2. Between 3 and 7 MHz the velocities are frequency independent

within this system's resolution capability.

3. Transit-time error, usually ascribed to bond and transducer

thickness, was found to be nonexistent.

4. Elastic constants of stainless steel 304 show surprisingly

little lot-to-lot variation, 0.9 to 1.4 percent depending on the

elastic constant. Larger variations reported in the literature are

due mainly to experimental measurement errors.

ACKNOWLEDGMENT

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

of Fusion Energy.

IGC
REFERENCES

1. R. W. Swindeman, W. J. MacAfee, and V. K. Sikka, in Reproducibility and Acccuracy of Mechanical Tests , STP 626 (Amer. Soc. Test. Mater., Phila.,

1977), pp. 41-64.

2. N. H. Fahey, Watertown Arsenal Laboratories Report WAL TR 118 (1960).

3. J. F. Haskins, General Dynamics Report ERR-AN-055 (1961).

4. P. E. Armstrong and D. T. Eash, in Advances in Cryogenic Engineering , Volume 14 (Plenum, New York, 1969), pp. 64-70.

5. H. M. Ledbetter, W. F. Weston, and E. R. Naimon, J. Appl . Phys., 3855-3860

(1975).

6. J. P. Hammond, M. W. Moyer, and C. R. Brinkman, in Proceedings Second

International Conference on Mechanical Behavior of Materials (Amer. Soc

.

Metals, Metals Park, Ohio, 1976), pp. 1037-1041.

7. J. F. Bates, Hanford Engineering Development Laboratory Report HEDL-SA-1068

(1976).

8. J. T. Richards, in Symposium on Determination of Elastic Constants , STP 129 (Amer. Soc. Test. Mater., Phila., 1952), pp. 71-100.

9. D. Lazarus, Phys. Rev. 7^, 545-553 (1949).

10. J. R. Neighbours, F. W. Bratten, and C. S. Smith, J. Appl. Phys, 2^, 389-393

(1952).

11. C. S. Smith and J. W. Burns, J. Appl. Phys. 15-18 ( 1953).

12. S. Eros and J. R. Reitz, J. Appl. Phys. 683-686 (1958).

13. L. D. Landau and E. M. Lifshitz, Theory of Elasticity (Pergamon, London,

1959), p. 13, p. 99.

14. M. Kikuchi, Trans. Jap. Inst. Met. j_2, 417-421 (1971 ).

187
TABLE

I.

Chemical

compositions

(weight

percent),

hardness,

and

mass

density.

>5'-'

4-roE , CO o o to to to OsJ 1 o o LO CO LO to LT)

t/^ 00 o in CM CO LO r— r--. r— CM CM to CT. cn 00 1

r—

CM 1

CO

c/0 c 00 C3^ CTi cr. 00 00 00 oo CTi CT) 00 CO CO 00 cri 00 CO 00 CTl ro <u CT! s; Q 1^ r--

CQ

to

to I

O) t—

c: O) o S

s-fC o m od

csjr^totoc\j'X)tDtoc\Jc\jcoLnLn'^cricor-^tOf^^ r^oor^cTitor^i^r^cocTir^r^r^r^r^r^i^r^cTir-'

' !)

i o

C_)

too

roo

I 3

^ I

CNJ

C\J I

un c CM ro

. . . ro

13

C_)

r-^ un CM

C» O I— CM

CM . . .

to CM CM

(C .— O

too

oo o

o

CO

f— I— CM I

o o o c o

o LO C3^ CM CM CM CM r— CO C\J O CM O O O o o CM

D. • O O

O 00

CM to to CO to 1 O CM O o o o o LD

• O o

oo

to CM

• CO

CT) 00 CO CO

LO CO CO to LO CO

>

I

’ \

) i

1

i

I

i

i

i- to

in CM CO CO ro

• to 00 ^ 00 1— CM

r— • • o o

i

00

1

1-^ LD LO CTi p— cn 00

• • r-- LO to to

cn o • 00 00 CTi oo CTi CTi

i o LO to to CM

00 CO LD to to LD

00 00 00 00 CO 00 CO 00 .4

' ' *“ ' ‘ ^ ' * o CM CO ^ LO to 00 CTi o CM 3 4 5 6 7 8 9 1 '

— r- — .— CM

i

I

i
Table II. Reproducibility and frequency effects for longitudinal and shear

velocities in specimen 5. Each entry is an average of 10 measurements.

Frequency

(MHz)

Echoes

(cm/ys)

Vt

(cm/ys)

3 1-2 .5762 .3127

2-3 .5763 .3126

3-4 .5761 .3126

5 1-2 .5752 .3123

2-3 .5754 .3123

3-4 .5757 .3124

7 1-2 .5750 .3126

2-3 .5751 .3127

3-4 .5752 .3128

Mean (1-2) .5755 ± .0006 .3125 ± .0002

Mean (1-3) .5756 ± .0006 .3125 ± .0002

Mean (3-4) .5757 ± .0005 .3126 ± .0002

Mean (3 MHz) .5762 ± .0001 .3126 ± .0001

Mean (5 MHz) .5754 ± .0003 .3123 ± .0001

Mean (7 MHz) .5751 ± .0001 .3127 ± .0001

Mean (overall

)

.5756 ± .0005 .3126 ± .0002

1^9
TABLE III. Effects of nonresonant driving frequency on measured ultrasonic

velocities in specimen 5.

Frequency

(MHz) (cm/ys)

Vt

(cm/ys)

2.5 .5759 a 3.0 .5761 .3127

3.5 .5758 .3125

4.0 .5756 .3126

4.5. .5756 .3126

5.0^ .5754 .3125

5.5 .5754 .3127

6.0 .5754 .3126

6.5 .5753 .3126

7.0 .5751 .3125

Mean .5755 .3128

Std. dev. .0003 .0005

^Irreproduci ble due to poor echo pattern.

^Fundamental frequency.

190
TABLE IV. Directional variation of ultrasonic wave velocities in specimen 5.

Propagation Polarization V.

Di rection Di rection (cm/ys) (cm/ys)

100 100 .5760

010 .3126

001 .3123

010 010 .5756

100 .3125

001 .3123

001 001 .5761

100 .3122

010 .3124

Mean .5759 .3124

Std. dev. .0003 .0001

Pet. std. error .5 .03

191
TABLE V. Ultrasonic velocities and elastic constants for 20 randomly selected

stainless-steel-304 samples.

Specimen (cm/ys)

V.

(cm/ys)

B (lO^^N/m^)

E (lO^^N/m^)

G

(lO^^N/m^) V

1 .5772 .3144 1.581 2.001 .776 .289

2 .5746 .3122 1.586 1.993 .772 .291

3 .5749 .3131 1.584 2.004 .777 .289

4 .5752 .3108 1.598 1.977 .764 .294

5 .5756 .3126 1.582 1.981 .767 .291

6 .5745 .3182 1.524 2.024 .791 .279

7 .5775 .3134 1.595 1.998 .774 .291

8 .5786 .3131 1.594 1.980 .766 .293

9 .5762 .3161 1.574 2.034 .791 .285

10 .5759 .3108 1.608 1.983 .766 .295

11 .5764 .3137 1.596 2.015 .781 .290

12 .5760 .3152 1.567 2.009 .781 .286

13 .5781 .3123 1.612 1.993 .770 .294

14 .5786 .3126 1.615 1.997 .772 .294

15 .5746 .3118 1 .581 1.979 .766 .291

16 .5761 .3131 1.592 2.002 .776 .290

17 .5752 .3125 1.568 1.970 .763 .291

18 .5746 .3129 1.570 1.985 .770 .289

19 .5736 .3138 1.545 1.980 .77u .286

20 .5740 .3147 1.567 2.020 .786 .285

Mean .5759 .3134 1.582 1 .996 .774 .290

Std. dev. .0015 .0018 .022 .017 .009 .004

Pet. std. error

0.3 0.6 1.4 0.9 1.1 1 .4

1

1

•

I

192
FIGURE CAPTIONS

Figure 1. Nine sound-velocity modes for an orthogonal system, where n_ denotes propagation vector and £ denotes polarization vector,

indicated by arrows in diagram.

Figure 2. Composite oscilloscope display showing pulse and first two

echoes. As described in text, transit-time measurements

were made between crests of leading cycles of consecutive

echoes

.

Figure 3 Schematic circuit diagram of ultrasonic pulse-echo velocitymeasurement system.

193
1

2

3

4

5

6

7

8

9

10 0 10 0

0 1 0

0 0 1

0 10 0 10

1 0 0

0 0 1 0 0 1 0 0 1

1 0 0

0 1 0

Figure 1. Nine sound-velocity modes for an orthogonal system, where n_ denotes propagation vector and £ denotes polarization vector,

indicated by arrows in diagram.
Figure 2. Composite oscilloscope display showing pulse and first two

echoes. As described in text, transit-time measurements

were made between crests of leading cycles of consecutive

echoes.

195
FREQUENCY COUNTER

u

S- +J i-

(_) cu

O) o I

CD "O O 00 N £Z >>

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"aj>o

x:0

<L)1<DCO

=3CLCow

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s- 4->

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scr> rO -o

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19C

measurement

system.
LOW TEMPERATURE SPECIFIC HEAT OF TWO STAINLESS STEELS

Wichita State University

Nati onal Bureau of Standards

197

Specific heat values between 2 and 20 K are reported for the two stainless steels AISI 304 and AISI 304L. Measurements were made by standard adiabatic calorimetry.

The results suggest that the thermal properties are roughly comparable to AISI 3WS,

but the absence of magnetic dusters implies that they are more appropriate for low

temperature work.

Low temperature specific heat of two stainless

steels

J.C. Ho, G.B. King and F.R. Fickett

Stainless steel is often used as a structural material in cryogenic systems. For many large-scale, low temperature

applications, knowledge of specific heat values for such

alloys is essential for proper design. Collings et al' have

recently made calorimetric measurements on AISI 3 1 OS

stainless steel. These measurements indicate the occurrence

of superparamagnetism, presumably associated with magnetic

clusters, by a temperature -independent contribution to the

specific heat. Corsan and Mitchem^ have recently rcjrorted

specific heat measurements from 4 K to room temperature

on a series of laboratory-produced Fe-Cr-Ni alloys with chronuum contents ranging from I 2-24 wt.A and nickel contents in the range I 2—20 wt.'/f. Differences of up to 307r were observed in the specific heats as measured at 4 K. However, to the best of our knowledge, no data previously

existed on the most common low temperature structural

stainless steel -AISI 304. This note describes the results of low temperature specific heat measurements i>n AlSi 304

and AISI 304L stainless steels. L designates low carbmi

content.

Experimental

The experimental materials were obtained from commercial

suppliers. The chemical analyses made on these specific

James C. Ho and Galen B. King are at the Physics Department,

Wichita State University, Wichita, Kansas 67208, USA. Fred R.

Fickett is at the Cryogenics Division, National Bureau of Standards,

Boulder, Colorado 80302, USA. The work was sponsored by the

Division of Magnetic F usion Energy of the Department of Energy. Received 15 February 1978,

materials are given in Table 1 . The specimens were macliined from the stock as received and were given no further heat treatments. Tlie specimens were cylinders 2.54 cm long and

2.22 cm in diameter with a mass of ~ 0.080 kg.

The specific heat measurements were made between 2 and

20 K by standard adiabatic calorimetry. The sample

temperature determinations were based on the electrical resistance of a commercial germanium thermometer.

Pulsed Joule heating was obtained by passing a current through a manganin wire which, along with the thermometer, was tliermally anchored to the sample. The heat capacity of the thermometer-heater assembly was separately measured

for addenda corrections. The performance of the calorimeter was checked by measuring the specific heat of a high-purity copper sample; the data were in good agreement with

literature values.

Results and discussion

The specific heal data are presented in Figs 1 and 2 in the formal of CJT versus T^, where C is the specific heat and

T \s tlic ab.solute temperature. The solid lines, represented

by the equations given on the figures are obtained by

least-squares fit with a standard deviation of 0.17%. The

linear term is due to the conduction electrons, while the others are lattice contributions.^ The fact that one needs more than just the term reflects the contribution ol higher-fiequency phonons at finite temperatures. The

overall uncertainty in the specific heat is estimated to be

Table 1. Chemical analysis of steels

Wt% of listed element

AISI No C Mn P S Si Cr Ni N Mo Co

304 0.061 0.820 0.020 0.017 0.590 18.30 8.45 - - -

304 L 0.030 1.430 0.025 0.013 0.031 18.50 8.75 0.061 0.29 0.18

001 1-2275/78/1805-0296 $01.00

199

CRYOGENICS . MAY 1978
Table 2. Heat capacity, entropy and enthalpy of AISI 304 and AISI 304L stainless steel

Temperature, Heat capacity. Entropy, Enthalpy, K J kg' 1 ^-1

J kg'-iK-i Jkg-’

304 304 L 304 304 L 304 304 L

2.0 0.936 0.843 0.0 0.0 0.0 0.0 3.0 1.41 1.27 0.469 0.422 1.17 1.06 4.0 1.89 1.70 0.939 0.847 2.82 2.54

5.0 2.37 2.15 1.41 1.27 4.95 4.47

6.0 2.87 2.60 1.88 1.70 7.57 6.84

7.0 3.37 3.07 2.37 2.14 10.7 9.67

8.0 3.89 3.56 2.85 2.58 14.3 13.0

9.0 4.43 4.06 3.34 3.03 18.5 16.8 10.0 4.99 4.59 3.84 3.48 23.2 21.1 n.o 5.57 5.14 4.34 3.95 28.5 26.0

12.0 6.18 5.73 4.85 4.42 34.4 31.4

13.0 6.82 •6.35 5.37 4.90 40.8 37.4

14.0 7.51 7.00 5.90 5.40 48.0 44.1 15.0 8.23 7.70 6.44 5.90 55.9 51.5

16.0 9.00 8.45 7.00 6.43 64.5 59.5

17.0 9.83 9.25 7.57 6.96 73.9 68.4

18.0 10.7 10.1 8.16 7.52 84.2 78.1 19.0 11.7 11.0 8.76 8.09 95.4 88.6 20.0 12.7 12.0 9.39 8.68 108 100

witliin 1%. Fig. 3 shows the temperature variation of the and AISI 304L are roughly comparable to AISI 31 OS.'

specific heat directly. However, the absence of magnetic clusters in these material

Based on the experimental specific heat equations, entrop> and enthalpy (heat content) values at temperatures up to 20 K have been calculated, with reference to 2 K, They

are summarized in Table 2. The results thus obtained suggest

that, as far as thermal properties are concerned, AISI 304

suggests that they are more appropriate in those low

temperature applications where magnetic properties are

critical. We have also calculated values of the Debye

temperature with the results: Oo = 496 K for 304 and

On = 457 K for 304L.

80 160 240 320 400 80 160 240 320 400

Fig. 1 C/T versus for AISI 304 Fig. 2 C/T versus for AISI 304L

CRYOGENICS . MAY 1978

200
Fig. 3 Specific heat of AISI 304 and AISI 304L as a function of temperature

It is also interesting to note that the data presented here

agree quite well with those for the 24Cr-12Ni alloy

measured by Corsan and Mitchem,^ which is the only one of their alloys with the same Ni : Cr ratio as AISI 304.

This seems to bear out their observation that it is this ratio which is correlated with the variation in specific heat among the alloys, a high ratio leading to a relatively high

specific heat.

References

1 CoUings, E.W., Jelinek, F.J., Ho, J.C., Mathui, M J*. Advances

in Cryogenic Engineering, Ed K.D. Tinimerhaus, R.P. Reed

and A. I'. Clark, 22 (Plenum Publishing Corp, NY 1977) 159

2 Corsan, J.M., Mitchen, N.l. Proc 6th Int Cryo Eng Conf

(IPC Sci and Tech Press, Guildford, Surrey, England 1976) 342

3 1 see eg, Kittel, C. Introduction to Solid State Physics, 4th Ed, (John Wiley and Sons. NY 1971) 252

4 Barron, T.H.K., .Monison, J.A. Can J Phys 35 (1957) 799

201

CRYOGENICS . MAY 1978
SPECIFIC

HEAT.

10^

Btu/(lb„

°F)

TEMPERATURE. K

4 6 8 10 20 30 40 60 80 100 200 300

1000

800

600

400

300

200

100

80

60

40

30

20

10

8

6

4

3

2

-450 -440 -420 -400 -300 -200 -100 0 80

TEMPERATURE.“F

Figure 4. This graph, from the NBS handbook LNG Materials and Fluids

(1977), shows how the data determined here for AISI 304 and

304L fit with specific heat data for other stainless steels.

202

SPECIFIC

HEAT.

J/(kg

K)
MAGNETIC SUSCEPTIBILITY AND MAGNETIZATION STUDIES

OF SOME COMMERCIAL AUSTENITIC STAINLESS STEELS

Battel le Co\umbus Laboratories

203

FOREWORD

Magnetic susceptibility measurements were carried

out at Battel le by R. D. Smith, and the magnetization

measurements were performed at The Ohio State

University by S. C. Ha'rt. The work was carried out

under the auspices of the Thermophysical Properties

Division, National Bureau of Standards, Boulder,

Colorado, F. R. Fickett, program monitor.

E. W. Collings

Principal Investigator

BATTELLE

BATTELLE — COLUMBUS

205
MAGNETIC SUSCEPTIBILITY AND MAGNETIZATION STUDIES

OF SOME COMMERCIAL AUSTENITIC STAINLESS STEEL ALLOYS

to

THE THERMOPHYSICAL PROPERTIES DIVISION

NATIONAL BUREAU OF STANDARDS

BOULDER, COLORADO 80302

INTRODUCTION

Magnetic susceptibility measurements, using the Curie magnetic

force technique, and magnetization measurements, using a vibrating-sample

magnetometer have been carried out on six AISI 300-series alloys 310S,

304, 304L, 304N, 316, 316L; as well as AWS 330 weld metal and Inconel

625, over the temperature ranges listed in Table 1.

TABLE 1. TEMPERATURE RANGES OF THE SUSCEPTIBILITY

AND MAGNETIZATION STUDIES

Alloy Name

Temperature Range of

Susceptibility Study (K)

Temperature Range of

Magnetization Study (K)

AISI 310S 77-414 4-75

AISI 304 9-408 4-143 plus 297

304L 77-410 4-112 plus 200, 296

304N 77-403 4-114 plus 294

AISI 316 77-407 3-110 plus 296

316L 77-413 4- 80 plus 181 , 286

AWS 330 77-416 4-65 plus 187, 295

In 625 5-412 --

After listing alloy compositions, describing sample preparation

procedures and so on, the numerical results of the study will be presented

BATTELLE — COLUMBUS

206
in two separate sections, the first dealing with the susceptibility results,

and the second with the magnetization. The report concludes with a summary

of the magnetic characteristics of austenitic stainless steels with

reference to the four principal types studied.

SPECIMEN MATERIALS

Sources and- Compositions

In Table 2 are listed the vendors and available specifications

of the alloys investigated. In Table 3, vendor-supplied compositions are listed together with average atomic weights (AW) calculated using the

expression AW = 100/E^. (x^./w^. ) > where x^. is the abundance in weight

percent of a constituent of atomic weight w^.

.

TABLE 2. VENDORS AND AVAILABLE SPECIFICATIONS

OF THE ALLOYS INVESTIGATED

Alloy Name

Supplier and Available

Specifications

AISI 31 OS Rolled Alloys, Inc., Heat No. 20,650

HRAP Plate, as-received

AISI 304 Armco Steel Corp., Heat No. 546,314

7/8-in. diameter bar, as-received

AISI 304L Jessop Steel Co., Heat No. 14,768

1-in. plate, received as 1900-1950 F (1038-1066 C), 1/2 h, water quench

AISI 304N Cameron Iron Works, Inc., CIW Serial No. 12585Z

Heat No. K2980, received as 1925 F (1052 C),

2-1/2 h, water quench

AISI 316 Commercial source 3/4-in. diameter bar,

as-recei ved

AISI 316L Armco Steel Corp., Heat No. 55,603, 1-in.

diameter bar, as-recei ved

AWS 330 Weld test pad, hence as-cast, rapidly chilled

from melt

BATTELLE — COLUMBUS

207
o

Li- o

tn2O

OQ

2:

O) cn 1

'I

O

O

: 1

rn O I •— — I 00 5 ' «=c

cc o c^.

a^ 00 LD ro ro 000

o

o

oo

o

o

oroO

O

COo

^ o

o o

COo

LO m CT^

CO CO cn

oro

CO

ro ro ro

CTi tno

CO

C\J r- o o

cr>

c\j r- ro

CO

c/^ 00

ci ct

o

d

«=j-

'!3

- O

<y^ LO

_ I:

208
Specimen Preparation

Samples were removed from the supplied material using a carborundum cutoff wheel, care being taken to prevent plastic (bulk)

damage during handling, and to remove surface damage by etching with

"glyceregia"*. The effect of bulk and surface damage on the magnetic

susceptibility of AISI 304N, for example, is discussed in Appendix A. Table 4 lists the degrees of etching administered, and the final

specimen weight prior to measurement of magnetic susceptibility and

magnetization, respectively. As can be seen from the tabl e , surface

TABLE 4. DEGREE OF ETCHING AND SPECIMEN WEIGHTS IN MAGNETIC

SUSCEPTIBILITY MEASUREMENT

Alloy Name

Amount

Removed

(%)

Fi na 1 Specimen

Weight (mg)

AISI BIOS 23 8.85

AISI 304 46 3.725

304L 40 4.99

304N 23 4.37

AISI 316 39 11.51

316L 33 11.50

AWS 330 17 4.95

In 625 -- 24.54

layers of thickness 0.010-0.015 cm were removed from the susceptibility

samples; after which an attempt was made to etch to about the same extent

the larger (1/3 to 1/2-g) magnetization samples.

* 10 ml HNO^, 20-50 ml HCl , 30 ml glycerol

.

BATTELLE — COLUMBUS

209
MAGNETIC SUSCEPTIBILITY STUDIES

Measuring Procedure

Magnetic susceptibilities were measured as functions of

temperature by the Curie technique using an electronic microbalance and

an electromagnet fitted with 18-cm diameter "constant-force" pole caps.

The vertical magnetic force field at the sample position (H-3H/8z) was calibrated at room temperature using a Vz-g sample of high-purity ^t

whose susceptibility and susceptibility temperature dependence are well

known. Ferromagnetic contamination of the specimen could be detected,

as well as corrected for, by the Honda-Owen method in which the matrix

susceptibility is obtained as the intercept, x„> of a reciprocal -field

plot based on the equation

X

= + M/H

in which M represents the magnetization (which must always be saturated)

of the ferromagnetic impurity.

If the parasitic ferromagnetic moment is zero or temperature

independent (as would be the case for Fe itself at temperatures below

about 500K*) the force-balance data are analysed by computer. On the

other hand if the reciprocal-field plots are non-linear, or if the slopes

turn out to be temperature dependent, the data are analyzed graphically.

Examples are given below.

Data Reduction

Graphical Analysis —AWS 330

Since AWS 330 showed evidence of superparamagnetism within the

liquid nitrogen temperature range (77-lOOK) the data were analysed by

* The saturated moment of Fe decreases by about 4% between 0 and 500K.

BATTELLE — COLLJMBUS

210
hand and presented as plots of x versus 1/H--Figures 1(a) and (b)

Since the plots seemed linear in the temperature range 129.5K-414.5K

extrapolated values of were acquired. These are presented in Table 5. Attempts were then made to computer fit the susceptibility temperature

dependence within this temperature range to the relationship

X

= C/(T-e)

Only by restricting the data to the temperature range above 172K could

an acceptable coefficient of variability (in this case 0.5%) be obtained.

The fitting parameters were then C = 11,214 ± 53 10’^cm^Kg~\ and

e = 111.8 ± 0.6 K, the latter value revealing the need for the reduced

temperature range, and indicating the advent of superparamagneti sm at

about that temperature, which presumably represents the Curie point of

pre-existing chemical clusters.

Hand Cal cul ation--AISI Type 304 (8.7-65.6K)

As indicated in Table 1, only for AISI 304 and In 625 were

force balance measurements extended into the liquid He range. In the case of case of AISI 304, since the susceptibility began to exhibit some field dependence in that range, the force balance data were reduced by

hand and presented under the headings "high" (10.0 kOe) and "low" (3.75

kOe) fields, as in Table 6, and plotted as in Figure 2.

Computer Analysis of Data, Mostly

Within the Temperature Range 77-400K

If inspection of the data indicated that the slopes of the

reciprocal -field plots would be either zero or temperature independent

they were analyzed by computer according to the algorithm illustrated

in Figure 3. An index of the data-sets analyzed in this way is given

in Table 7. The susceptibility temperature dependences themselves are presented numerically in Tables 8 through 14, and graphically for several

representative alloys in Figure 4.

BATTELLE — COLUCV1BUS

211
140 —

130 —

cn

Z5E

CD CO

IO

><

L5

"q. CD oCO

Z5 ID

CO

COo

o

o3

!Zcmo

120

NO100

90

80

70

60

50

40

30

1 r

-X—X- -X-

-X—X- -XX- —

X—X.

X ^X —

—X—X ~~

N/ X- —

N/ ^x -X—

X

N/ ^x V X

N/ X —

XX NX

XN X

X — NX XX

^ X ^

. . . NX NX X—X -X X X

196.5 K

209.0 K -

2230 K

233.5 K

246.5 K

258.5 K

271.5 K

2800 K

293.0 K

3065 K

324.5 K

347.0 K

370.0 K

391.5 K

414.5 K

0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.4 2.6 2.8 3.0 3.2 3.4 3.6 3.8 4.0

Reciprocal Magnetic Field, H~', lO”"^ 0e~'

FIGURE 1(a). MAGNETIC SUSCEPTIBILITY OF AWS 330. RECIPROCAL FIELD

PLOTS FOR THE TEMPERATURE RANGE 41 4. 5-1 96. 5K. THE GRADUAL

DEVELOPMENT OF "FERROMAGNETIC SLOPE" IS EVIDENT.

212
Magnetic

Mass

Susceptibility,

x,

lO^emu

i

I

i

FIGURE 1(b). MAGNETIC SUSCEPTIBILITY OF AWS 330--C0NTINUED . RECIPROCAL

FIELD PLOTS FOR THE TEMPERATURE RANGE 185-77K. THE ONSET

OF SUPERPARAMAGNETISM AT TEMPERATURES NEAR TOOK IS EVIDENT.

213
TABLE 5. MAGNETIC SUSCEPTIBILITY OF AWS 330. GRAPHICALLY

EXTRAPOLATED X VALUES FOR THE TEMPERATURE RANGE

414. 5-129. 5K "

Temperature

(K)

Magnetic

Susceptibil ity

(10"“cm39"^

)

Temperature

(K)

Magnetic

Susceptibility

(10“°cm^g"

414.5 37.8 258.5 77.0

391.5 40.6 246.5 83.5

370.0 43.8 233.5 92.4

347.0 47.0 223.0 101.1

324.5 52.5 209.0 116.0

306.5 57.5 196.5 132.7

297.0 60.1 185.0 155.3

296.5 60.2 171.5 185.7

156.5 233.5

293.0 61.5

280.0 65.9 142.5 303.3

272.0 69.1 129.5 383.0

BATTELLE — COLUMBUS

214
TABLE 6. MAGNETIC SUSCEPTIBILITY OF AISI TYPE 304 IN

THE TEMPERATURE RANGE 8.7-65. 6K

Magnetic Field, 10.0 kOe Magnetic Field, 3.75 kOe

Temperature

(K)

Magnetic

Susceptibility

(10"6cm3g"^

)

Temperature

(K)

Magnetic

Susceptibility

(10"°cm3g-l

)

8.7 69.8 8.9 71 .7

9.3 69.9 9.5 72.1

9.4 69.9 9.5 71 .3

10.6 70.1 10.7 72.2

11.6 70.3 11.5 71.3

12.3 70.4 12.4 72.4

13.7 70.8 14.1 72.9

15.4 71 .0 15.2 72.3

17.2 71.6 16.9 72.8

18.9 72.1 18.5 73.2

20.7 72.7 20.2 74.1

24.1 74.3 23.7 75.8

26.1 75.2 25.6 76.2

31.4 78.1 30.1 78.2

34.5 81 .2 33.5 81 .7

38.1 84.4 37.5 83.7

40.0 85.2 40.0 87.8

45.2 82.8 45.1 86.4

51.7 75.6 50.7 78.6

64.7 63.6 65.6 64.3

BATTELLE — COLUMBUS

215
CDO

Ql

E

,CD

|_E> X ‘A|i|iqi|d03sn2 oijauboiAj

216

FIGURE

2.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCE

OF

AISI

TYPE

304,

AT

FIELDS

LEVELS,

WITHIN

THE

TEMPERATURE

RANGE

8.7-65.6K.
± std dev X

i

I

Reject all Xjnt's >3 std devs

from mean, retain same

average slope, re-compute

intercepts and std devs for

each T

Tabulate

T, X ± std dev x

For each T fit straight

line through X versus

H’' data

Put lines of equal slope

(=(slope>n) through ajj

datum points. For each

set of five (each T) tabulate <Xint >5 ± std dev X

3 . ALGORITHM FOR THE COMPUTER REDUCTION OF SUSCEPTIBILITY

FORCE BALANCE DATA
TABLE 7. INDEX OF COMPUTER-PRODUCED

TEMPERATURE DEPENDENCE DATA

SUSCEPTIBILITY

TABLES

Temperature Range

of Computer Analysis Table

Alloy Name (K) Number

AISI 310S 77-414 8

AISI 304 77-408 9

304L 77-410 10

304N 77-403 11

AISI 316 77-407 12

316L 77-413 13

In 625 9-80* 14(a)

77-412 14(b)

* Results normalized at 77K to those of the higher

temperature run.

BATTELLE — COLUMBUS

218
TABLE 8. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

AISI 310S OVER THE TEMPERATURE RANGE 77-414K

Temperature

(K)

Susceptibility

at

Intercept

Number of

Poi nts

Accepted

Standard Deviation

of

Intercept

77.0 86.015 4 0.053

87.0 76.787 3 0.142

85.0 77.670 3 0.132

96.5 67.705 4 0.095

106.0 61 .112 4 0.099

118.5 54.540 4 0.030

129.5 49.934 5 0.081

142.5 45.891 5 0.079

153.0 42.927 5 0.055

166.0 40.057 5 0.056

177.5 37.891 5 0.051

189.0 36.056 5 0.053

202.5 34.278 5 0.061

217.5 32.433 5 0.085

231.5 31.072 5 0.057

245.0 29.975 5 0.057

257.0 28.964 5 0.048

271.0 28.002 5 0.046

281.5 27.335 5 0.036

293.5 26.605 5 0.070

298.5 26.241 5 0.084

305.5 25.990 5 0.069

326.5 24.918 5 0.064

347.5 24.081 5 0.050

371.0 23.196 5 0.042

391 .0 22.523 5 0.079

414.0 21.865 5 0.040

BATTELLE — COLUMBUS

219
TABLE 9. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

AISI 304 OVER THE TEMPERATURE RANGE 77-408K

Temperature

(K)

Susceptibility

at

Intercept

Number of

Points

Accepted

Standard Deviation

of

Intercept

77.0 55.349 4 0.066

93.5 48.675 4 0.036

103.5 45.610 4 0.034

118.5 42.033 4 0.106

131.5 39.060 5 0.074

143.5 37.598 5 0.070

152.5 36.046 5 0.046

166.5 34.870 5 0.057

182.0 33.430 5 0.084

194.0 32.286 5 0.092

203.0 31.599 5 0.144

213.5 30.897 5 0.087

228.5 30.125 4 0.092

241 .0 29.401 5 0.039

251.0 28.851 5 0.082

265.5 28.165 5 0.105

277.5 27.710 5 0.076

289.0 27.221 5 0.103

300.5 26.930 5 0.092

302.0 26.786 5 0.090

322.5 26.095 5 0.099

344.0 25.519 5 0.063

366.0 24.939 5 0.133

388.0 24.421 5 0.064

408.0 24.004 5 0.076

BATTELLE — COLUMBUS

220
TABLE 10. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

AISI 304L OVER THE TEMPERATURE RANGE 77-41 OK

Temperature

(K)

Susceptibil ity

at

Intercept

Number of

Points

Accepted

Standard Deviation

of

Intercept

77.0 50.460 4 0.041

94.0 44.752 3 0.031

107.0 41.268 . 5 0.070

125.0 38.038 5 0.077

141.0 35.625 5 0.078

158.0 33.698 4 0.027

172.5 32.328 5 0.042

189.0 31 .106 4 0.048

202.0 30.266 5 0.063

213.0 29.527 5 0.057

223.5 28.972 5 0.073

236.0 28.277 5 0.041

248.5 27.692 5 0.050

261 .0 27.232 5 0.064

274.0 26.629 5 0.067

287.0 26.181 5 0.037

298.5 25.888 5 0.078

302.0 25.738 5 0.083

321.5 25.065 5 0.049

343.5 24.411 4 0.062

367.0 23.848 5 0.083

387.5 23.362 5 0.043

409.5 22.954 4 0.060

BATTEULE — COLUCVIBUS

221
TABLE 11. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

AISI 304N OVER THE TEMPERATURE RANGE 77-403K

Temperature

(K)

Susceptibil ity

at

Intercept

Number of

Points

Accepted

Standard Deviation

of

Intercept

77.0 52.388 5 0.084

90.5 47.291 5 0.061

104.5 43.122 4 0.033

116.0 40.438 5 0.065

126.5 38.366 5 0.090

138.0 36.600 5 0.043

152.5 34.814 5 0.046

168.0 33.103 5 0.063

183.5 31 .801 4 0.062

197.0 30.747 5 0.086

211.5 29.765 5 0.093

226.0 28.915 5 0.064

239.5 28.152 5 0.107

251 .5 27.651 5 0.082

266.5 27.019 4 0.098

280.5 26.527 4 0.072

296.5 25.972 5 0.064

298.5 26.004 5 0.040

298.5 25.958 5 0.065

312.5 25.550 5 0.041

336.5 24.840 4 0.056

358.0 24.261 5 0.080

380.5 23.787 5 0.079

402.5 23.406 5 0.087

BATTELLE — COLUIVIBUS

222
TABLE 12. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

AISI 316 OVER THE TEMPERATURE RANGE 77-407K

Temperature

(K)

Susceptibil ity

at

Intercept

Number of

Points

Accepted

Standard Deviation

of

Intercept

77.0 75.446 4 0.030

96.5 62.727 4 0.034

110.5 55.576 - 4 0.082

124.0 50.688 5 0.042

136.5 47.005 5 0.031

149.5 43.967 5 0.045

164.0 41.348 5 0.044

178.0 39.143 5 0.046

189.0 37.706 5 0.058

198.5 36.602 5 0.065

209.5 35.461 5 0.050

222.0 34.304 5 0.034

233.0 33.366 5 0.049

243.0 32.608 5 0.063

254.5 31 .838 5 0.053

264.0 31.250 5 0.041

278.0 30.425 5 0.047

291.5 29.753 5 0.042

300.0 29.345 5 0.054

306.5 28.980 5 0.059

319.0 28.446 5 0.065

344.0 27.462 4 0.056

366.0 26.683 5 0.042

386.5 26.051 5 0.053

407.0 25.482 5 0.061

BATTELCE — COLUIVIBUS

223
TABLE 13. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

AISI 31 6L OVER THE TEMPERATURE RANGE 77-41 3K

Temperature

(K)

Susceptibil ity

at

Intercept

Number of

Points

Accepted

Standard Deviation

of

Intercept

77.0 78.875 5 0.085

89.5 69.635 3 0.073

102.5 61.671 5 0.057

116.0 55.031 5 0.060

127.0 51 .395 4 0.026

137.0 48.159 4 0.016

148.0 45.371 5 0.030

160.0 42.993 5 0.054

169.5 41.313 5 0.042

179.0 39.785 5 0.041

190.0 38.301 4 0.054

203.0 36.869 5 0.038

214.5 35.666 5 0.046

224.5 34.733 5 0.043

234.0 33.857 5 0.039

247.0 32.883 5 0.030

258.0 32.145 4 0.013

270.5 31.327 5 0.060

281.5 30.666 5 0.051

295.0 29.958 5 0.059

297.0 29.766 5 0.042

309.0 29.275 5 0.034

324.0 28.583 5 0.034

346.0 27.755 5 0.047

368.5 26.936 5 0.053

390.0 26.236 5 0.052

413.0 25.582 5 0.038

BATTELLE — COLUMBUS

224
TABLE 14(a). MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE

OF INCONEL 625 OVER THE TEMPERATURE RANGE 9-80K

Temperature

(K)

Susceptibility

at

^ InterceptNumber of

Poi nts

Accepted

Standard Deviation

of

Intercept

9.5 33.923 5 0.065

9.5 34.170 3 0.087

10.1 32.564 4 0.081

10.5 31.774 4 0.069

11.3 29.971 5 0.095

12.5 28.058 5 0.102

13.4 26.629 5 0.037

15.8 23.718 5 0.030

19.9 20.287 5 0.066

23.8 17.904 5 0.016

29.4 15.481 5 0.074

34.3 13.692 5 0.044

38.2 12.740 4 0.053

42.8 11.806 5 0.077

48.7 10.897 5 0.055

56.0 10.083 5 0.054

61 .3 9.490 5 0.035

62.5 9.412 5 0.062

68.1 8.995 5 0.042

74.2 8.510 5 0.039

79.6 8.249 4 0.056

* After normalization, at 77K, to the data of Table 14(b) (normalization factor: 1.1152).

BATTELLE — COUUtVlBUS

225
TABLE 14(b). MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCE OF

INCONEL 625 OVER THE TEMPERATURE RANGE 77-41 2K

Temperature

(K)

Susceptibil i ty

at

Intercept

Number of

Points

Accepted

Standard Deviation

of

Intercept

77.0 8.354 5 0.035

90.0 7.784 5 0.027

97.0 7.540 5 0.018

98.5 7.466 5 0.019

110.0 7.083 4 0.022

110.0 7.091 4 0.011

121 .0 6.790 5 0.024

122.0 6.761 5 0.020

131.5 6.557 5 0.026

141 .5 6.363 5 0.028

153.0 6.134 5 0.034

161 .0 6.014 5 0.033

170.0 5.874 5 0.031

180.0 5.746 5 0.054

194.0 5.608 5 0.024

203.0 5.506 5 0.039

211.5 5.450 4 0.018

226.0 5.294 4 0.033

233.5 5.236 5 0.047

242.5 5.189 5 0.037

250.5 5.139 5 0.033

259.5 5.096 4 0.023

270.0 5.032 4 0.027

282.0 4.961 5 0.039

294.0 4.910 4 0.019

296.5 4.916 5 0.026

320.0 4.792 5 0.028

343.0 4.698 5 0.037

366.0 4.616 5 0.023

388.0 4.559 4 0.017

412.0 4.479 5 0.023

BATTELLE — COLUMBUS

226
FIGURE 4. MAGNETIC SUSCEPTIBILITY TEMPERATURE DEPENDENCES FOR AISI 310 S,

AISI 304, AISI 316, AWS 330, AND In 625

227
Curve Fitting

With the aid of the well-known SPSS (1967) package the data of

Tables 8 through 14 were fitted to the relationship

X = Xq + C/(T-e)

In addition, as has been mentioned, those of Table 5 (AWS 330) were fitted

to

Usually three iterations were employed in order to improve the goodness

of fit (expressed in terms of the coefficient of variability) by reducing

the temperature range and eliminating what appeared to be bad data. The

coefficients of the equations are listed in Table 15. Also listed in the

table is an average (over all atoms in the alloy) effective moment per

atom, p, in Bohr magnetons, calculated from the relationship

where C is the so-called "Curie constant".

* Statistical Package for the Social Sciences, Northwestern University,

Version 6.5, April 1, 1976.

X

= C/(T-0)

2/8
TABLE

15.

COEFFICIENTS

OF

THE

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

SOME

300-SERIES

STAINLESS

STEELS

AND

INCONEL

625

CJ

•r“ Q. CO

(U

C LD ro o ro CO LO

C •'

O o r— ro (

1 o O) CD E +-> o O O O o o o 03 O OJ • • • • • • • 03 ^ -M C

S- 03 CD o o o +1 +1 +1 o o+1 -H 1 1

<D O) \ 03 1

ro CM 1— 1— CM CD > > +-> E CO 00 CM 1— ro CM

ea; •!— c n f— r— r— ro ro CM

4-> CU S- • • • • • • • o E .C 0)0 0

4- s: CQ

<4-

CM

UJ

4->

11 OJ CD >, C03 O LO ro CM r— CO H 1 1 1 03 1—

SO

03 00

O CD CO CO 1 1 0) •<- ca: LO ^ LO LO LO 1 > E

cc

O 4-

LT> LO LO LO LO to LO

-M

O c -

cn CT) C\J 1— CO ro r-^ CO LO ro

CO

-H 1— LO

+1 +1 +1 <51 ‘Tl'Ti o E oo c— CM ro ca- CM ro 00 ^ o E CD 00 CO CO CD r— LO LO CO CD O O 1— CO 00 CD CM ro 1 1 CO ro CM ro ro r\ o o 1 1— f—

CD 00 o ro CO 1— ro • o • • ^ Ti +1 ? ^i-Ti

CD lO +1 CO CO CO CO 00 ro CM • • CM • • ro • CO CM ro ca- f— CM «aC\J C30 >— 1 r— t— 1 + 1 1 + + ' 1 1 CO 1^ ^ 1 1 1 O o o o 03 CJ> • 1— CM • f— • • o E CO X o E CM o CD CM

CO CD CO CO 00 00 ca- o o CO LO CO LO LO CO CO r— r— 1 f— t—“ r— r—

«s ' •

>,

C 4-J O) -r- •r“U -r- ^ •r- C|_ J21 CM ro CM CM CM LO LO CM

c+-

O C4- -rO o o o O

O o OO

O) s- O 03

C_) > '

O-O E 00

S- =3 +-> O) 4-> C LO ca- CD o ca- CO 00 CM

O JD 03 -I— E Q O

13 Du

-Z.

CSJ CM 1— CM CM CM CM ro 1 O)

SZ3 -M OJ

O ro o LO I

C\J 03 CD^-' f— O CO o O CD t— CD 1— ! S- C

O) 03 ^ 1 1 1 1 ^ ro 1 1 i CLoa LO ca- 1

CO CM CO E 00 CD 1

•

• , O) f— ( 1—

oo _l z _i 11 >» o cd" ca- ca- CO CO o <u — o o o f— o E ro ro ro ro ro ro ro LO — ro CO

ct z: »—« »—«

CO i oo oo CO CO CO

—1

—t-H 3: cz <C ca; ca: ca: c *—•

229
MAGNETIZATION STUDIES

Measuring Procedure

Magnetization was measured as a function of field and temperature

using a vibrating-sample magnetometer. Prior to the commencement of each

temperature run, which lasted about eight to ten hours, the magnetometer

was calibrated and aligned using an NBS-traceabl e sample of pure Ni

.

Magnetization Results

An index of the temperature and field ranges covered in this

study is given in Table 16, while the results themselves are presented

numerically in Table 17 through 23 and graphically in Figures 5 through

11 .

TABLE 16. INDEX OF THE TEMPERATURE AND MAGNETIC

FIELD RANGES COVERED

Temperature Magnetic Field

Alloy Name Range (K) Range (kOe)

AISI 31 OS 4-75 0.5-13.33

AISI 304 4-143 plus 297 0.5-13.20

304L 4-112 plus 200, 296 0.05-13.00

304N 4-114 plus 294 0.01-13.00

AISI 316 3-110 plus 296 0.5-13.33

31 6L 4- 80 plus 181, 286 0.05-13.00

AWS 330 4- 65 plus 187,295 0.01-13.20

230
TABLE 17. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AISI TYPE BIOS

Samp !• 310--S Ma,gnet. i xat.i on V. e. Tempena'tLjr'e v. m. Appl led F 1 o 1 d. Un i -t.s “omu/ 0

i~am pa9o 1

H

T 4. 2 7. 0 10. 0 15. 0 20. 0 25. 0 30. 0 35. 0

500 0. 717 0. 711 0. 637 0. 444 0. 325 0. 236 0. 175 0. 136

1000 1. 256 1. 246 1 . 106 0. 833 0- 629 0. 463 0. 341 0. 267

nil 1. 362 1. 349 1. 195 0. 912 0. 693 0. 510 0. 378 0. 295

1250 1. 492 1. 473 1. 307 1 . 010 0. 775 0. 571 0. 426 0. 333

1429 1. 634 1. 613 1. 433 1. 124 0. 870 0. 648 0. 481 0. 376

1667 1 . 822 1 . 812 1 . 600 1. 277 1 . 002 0. 754 0. 558 0. 436

2000 2. 060 2. 026 1. 804 1. 468 1. 169 0. 883 0. 664 0. 521

2500 2. 372 2. 330 2. 089 1. 735 1. 407 1 . 082 0. 820 0. 645

3333 2. 814 2. 761 2. 497 2. 124 1. 767 1. 396 1. 076 0. 852

5000 3. 499 3. 444 3. 155 2. 766 2. 386 1. 955 1. 547 1. 248

1 0000 4. 893 4. 840 4. 549 4. 179 3. 756 3. 261 2. 743 2. 309

13333 5. 554 5. 501 5. 210 4. 840 4. 443 3. 967 3. 393 2. 923

3 -1 Units are: Magnetization, cm g ; Temperature, K; Magnetic field, Oe,

231
TABLE 17 continued

lamp 1 o 310‘S Magrlot 1 zat 1 on V. «. Tompor atuno V. e. App 1 lod Fiold. Unite' pago 2

H

j 40. 0 45. 0 50. 0 55. 0 60. 0 65. 0 70. 0 75. 0

500 0. 109 0. 092 0. 078 0. 068 0. 060 0. 055 0. 049 0. 045

1000 0. 215 0. 181 0. 154 0. 135 0. 120 0. 108 0. 098 0. 089

nil 0. 239 0. 202 0. 172 0. 150 0. 134 0. 120 0. 108 0. 099

1250 0. 269 0. 227 0. 193 0. 171 0. 151 0. 135 0. 123 0. 112

1429 0. 305 0. 257 0. 220 0. 192 0. 171 0. 153 0. 139 0. 127

1667 0. 355 0. 303 0. 257 0. 225 0. 199 0. 185 0. 163 0. 148

2000 0. 421 0. 358 0. 308 0. 268 0. 238 0. 213 0. 194 0. 177

2500 0. 524 0. 444 0. 381 0. 334 0. 296 0. 267 0. 241 0. 221

3333 0. 701 0. 590 0. 505 0. 444 0. 391 0. 354 0. 321 0. 294

5000 1. 026 0. 875 0. 756 0. 664 0. 587 0. 529 0. 479 0. 439

10000 1. 957 1. 701 1. 484 1. 312 1. 166 1. 053 0. 955 0. 878

13333 2. 528 2. 211 1. 947 1. 724 1. 547 1. 399 1. 272 1. 172

232
TABLE 18. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AISI 304

SaiTip 1 e T“304 Magnot i zat 1 on V. M. Tampei~a't.ur« v. «. Appl i ed Field. UniC^ “etnu/ gram page 1

H r 4.

2

7. 0 10. 0 15. 0 20. 0 25. 0 30. 0 35. 0

500 0. 036 0- 037 0. 039 0. 039 0. 039 0. 040 0. 041 0. 043

667 0. 051 0. 052 0. 050 0. 051 0. 051 0. 052 0. 055 0. 056

1000 0. 071 0. 072 0. 073 0. 074 0. 074 0. 077 0. 079 0. 083

nil 0. 079 0. 080 0. 001 0. 081 0. 083 0. 084 0. 088 0. 092

1250 0. 008 0. 009 0. 090 0. 091 0. 092 0. 094 0. 098 0. 102

1429 0. 099 0. 101 0. 102 0. 103 0. 105 0. 107 0. Ill 0. 116

1667 0. 117 0. 117 0. 118 0. 119 0. 122 0. 125 0. 129 0. 135

2000 0. 137 0. 139 0. 140 0. 141 0. 145 0. 148 0. 153 0. 160

2500 0. 170 0. 172 0. 173 0. 175 0. 179 0. 183 0. 189 0. 198

3333 0. 225 0. 227 0. 229 0. 231 0. 236 0. 242 0. 251 0. 263

5000 0. 335 0. 337 0. 340 0. 344 0. 352 0. 360 0. 374 0. 392

7500 0. 490 0. 502 0. 505 0. 511 0. 523 0. 536 0. 558 0. 586

10000 0. 664 0. 670 0. 673 0. 682 0. 698 0. 716 0. 748 0. 782

12500 0. 629 0. 838 0. 841 0. 857 0. 875 0. 900 0. 941 0. 984

13200 0. 070 0. 085 0. 891 0. 907 0. 925 0. 953 0. 997 1. 044

n 1 Units are: Magnetization, cm'^g ; Temperature, K; Magnetic field, Oe.

233
TABLE 18 continued

iumpl e r “304 Magnet i zat 1 on V. a. Temperatone v. e. AppI led Field. Unite"emo/gram page 2

H T 30. 0 41. 0 44. 0 47. 0 50. 0 55. 0 60. 0 70. 0

500 0. 046 0. 047 0. 045 0. 044 0. 042 0. 039 0. 036 0. 032

667 0. 060 0. 061 0. 059 0. 050 0. 055 0. 052 0. 047 0. 041

1000 0. 007 0. 009 0. 007 0. 004 0. 002 0. 075 0. 060 0. 060

nil 0. 096 0. 090 0. 097 0. 094 0. 091 0. 003 0. 076 0. 067

1250 0. 106 0. 110 0. 100 0. 107 0. 102 0. 093 0. 005 0. 075

1429 0. 119 0. 124 0. 123 0. 121 0. 116 0. 105 0. 097 0. 004

1667 0. 139 0. 144 0. 143 0. 139 0. 134 0. 123 0. 112 0. 100

2000 0. 167 0. 170 0. 170 0. 164 0. 159 0. 145 0. 134 0. 110

2500 0. 207 0. 211 0. 211 0. 204 0. 190 0. 101 0. 166 0. 147

3333 0. 273 0. 279 0. 201 0. 270 0. 261 0. 239 0. 220 0. 195

5000 0. 406 0. 414 0. 415 0. 402 0. 307 0. 355 0. 327 0. 902

7500 0. 604 0. 617 0. 614 0. 595 0. 573 0. 526 0. 406 0. 430

1 0000 0. 007 0. 019 0. 013 0. 700 0. 757 0. 692 0. 642 0. 573

12500 1. 012 1. 019 1. 009 0. 970 0. 920 0. 072 0. 004 0. 713

13200 1. 069 1. 075 1. 065 1. 031 0. 901 0. 922 0. 050 0. 757

234
TABLE 18 continued

Saitipl* T~304 Magnet 1 zat. i on v. a. Tamporat.ur« v. a. Appllad FlalcJ. (Jnit.a**aimj/ 0

rani paga 3

H T 100.0 143. 0 296. 5

500 0. 020 0. 021 0. 017

667 0. 037 0. 027 0. 023

1000 0. 051 0. 039 0. 032

nil 0. 055 0. 043 0. 036

1250 0. 063 0. 049 0. 039

1429 0. 071 0. 055 0. 044

1667 0. 082 0. 064 0. 051

2000 0. 097 0. 076 0. 059

2500 0. 116 0. 094 0. 073

3333 0. 156 0. 124 0. 096

5000 0. 230 0. 104 0. 141

7500 0. 341 0. 275 0. 209

10000 0. 449 0. 369 0. 270

12500 0. 564 0. 467 0. 340

13200 0. 595 0. 492 0. 369

235
TABLE 19. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AISI 304L

Sampl* 304--L Magn«t. 1 zat. 1 on V. •. T«mp«rat.ur« v. m. Appl l«d Fl*ld. Unlt.a“•mu/^ram pay* 1 H T 4. 2 15. 0 25. 0 30. 0 40. 0 50. 0 60. 0 70. 0

50 0. 0214 0. 0236 0. 0237 0. 0242 0. 0242 0. 0232 0. 0236 0. 0233

100 0. 0409 0. 0430 0. 0437 0. 0437 0. 0436 0. 0420 0. 0422 0. 0415

200 0. 0741 0. 0760 0. 0764 0. 0768 0. 0765 0. 0740 0. 0743 0. 0735

500 0. 166 0. 168 0. 170 0. 171 0. 170 0. 166 0. 166 0. 166

667 0. 211 0. 215 0. 216 0. 216 0. 216 0. 212 0. 214 0. 211

1000 0. 298 0. 298 0. 298 0. 301 0. 301 0. 299 0. 299 0. 294

1111 0. 320 0. 322 0. 322 0. 323 0. 323 0. 323 0. 321 0. 316

1250 0. 345 0. 348 0. 348 0. 349 0. 349 0. 348 0. 345 0. 342

1429 0. 370 0. 375 0. 376 0. 377 0. 377 0. 376 0. 376 0. 370

1667 0. 400 0. 408 0. 408 0. 408 0. 410 0. 411 0. 410 0. 403

2000 0. 443 0. 448 0. 449 0. 450 0. 450 0. 454 0. 455 0. 444

2500 0. 493 0. 499 0. 502 0. 502 0. 505 0. 509 0. 511 0. 497

3333 0. 563 0. 569 0. 571 0. 574 0. 578 0. 586 0. 589 0. 570

5000 0. 666 0. 675 0. 679 0. 682 0. 691 0. 707 0. 710 0. 683

7500 0. 790 0. 802 0. 808 0. 814 0. 830 0. 857 0. 860 0. 820

10000 0. 908 0. 920 0. 929 0. 938 0. 962 1. 001 1. 003 0. 949

12500 1. 021 1. 039 1. 053 1. 064 1. 092 1. 144 1. 142 1. 079

13000 1. 045 1. 062 1. 079 1. 087 1. 118 1. 172 1. 170 1. 104

-1 Units are: Magnetization, cm g” ; Temperature, K; Magnetic field, Oe.

23G
TABLE 19 continued

Somplo 304“L Magnet 1 zat 1 on v. e. Temperature v. o. Applied Field. Uni te**emu/gram page 2

H T 88. 0 112. 0 200. 0 295. 5

50 0. 0236 0. 0234 0. 0254 0. 0281

100 0. 0417 0. 0417 0. 0459 0. 0519

200 0. 0616 0. 0736 0. 0813 0. 0921

500 0. 165 0. 164 0. 175 0. 187

667 0. 209 0. 206 0. 215 0. 226

1000 0. 292 0. 287 0. 284 0. 288

nil 0. 314 0. 309 0. 303 0. 305

1250 0. 338 0. 330 0. 322 0. 322

1429 0. 362 0. 355 0. 344 0. 343

1667 0. 392 0. 384 0. 370 0. 367

2000 0. 432 0. 421 0. 403 0. 398

2500 0. 482 0. 467 0. 443 0. 436

3333 0. 550 0. 530 0. 495 0. 484

5000 0. 650 0. 620 0. 567 0. 549

7500 0. 770 0. 725 0. 649 0. 622

i 0000 0. 882 0. 824 0. 726 0. 691

12500 0. 993 0. 920 0. 799 0. 752

13000 1. 015 0. 940 0. 814 0. 769

237
TABLE 20. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AISI 304N

iampl«

H

304-“N Magnot. i zat. 1 on v. •. T«mp«rat.ijr« v. •. Appl 1 «d F

1

aid. Unl'ta** emu/gram paga 1

T 4. 2 10. 0 15. 0 20. 0 25. 0 30. 0 35. 0 40. 0

10 0. 0006 0. 0006 0. 0005 0. 0005 0. 0006 0. 0006 0. 0005 0. 0006

20 0. 0012 0. 0010 0. 0010 0. 0011 0. 0011 0. 0011 0. 0011 0. 001

1

30 0. 0015 0. 0016 0. 0015 0. 0016 0. 0016 0. 0017 0. 0017 0. 0017

50 0. 0028 0. 0025 0. 0026 0. 0027 0. 0026 0. 0027 0. 0028 0. 0028

100 0. 0050 0. 0051 0. 0051 0. 0052 0. 0053 0. 0054 0. 0055 0. 0057

200 0. 0112 0. 0101 0. 0102 0. 0104 0. 0106 0. 0108 0. 0111 0. 0115

300 0. 0161 0. 0154 0. 0154 0. 0156 0. 0161 0. 0164 0. 0166 0. 0173

500 0. 0250 0. 0252 0. 0256 0. 0260 0. 0263 0. 0271 0. 0276 0. 0287

667 0. 0329 0. 0335 0. 0340 0. 0344 0. 0352 0. 0358 0. 0370 0. 0305

1000 0. 0499 0. 0512 0. 0517 0. 0525 0. 0535 0. 0548 0. 0562 0. 0583

nil 0. 0557 0. 0567 0. 0573 0. 0586 0. 0592 0. 0609 0. 0628 0. 0648

1250 0. 0622 0. 0638 0. 0644 0. 0655 0. 0667 0. 0683 0. 0703 0. 0729

1429 0. 0708 0. 0725 0. 0733 0. 0746 0. 0758 0. 0778 0. 0802 0. 0830

1667 0. 0823 0. 0843 0. 0854 0. 0868 0. 0882 0. 0909 0. 0933 0. 096b

2000 0. 0983 0. 101 0. 102 0. 104 0. 106 0. 109 0. 112 0. 116

2500 0. 122 0. 126 0. 127 0. 130 0. 132 0. 135 0. 139 0. 144

3333 0. 163 0. 167 0. 169 0. 171 0. 175 0. 179 0. 185 0. 192

5000 0. 242 0. 250 0. 252 0. 258 0. 262 0. 270 0. 277 0. 280

7500 0. 365 0. 375 0. 380 0. 387 0. 394 0. 405 0. 417 0. 434

10000 0. 488 0. 501 0. 508 0. 518 0. 527 0. 541 0. 557 0. 581

12500 0. 613 0. 630 0. 639 0. 649 0. 662 0. 680 0. 701 0. 732

13000 0. 638 0. 654 0. 663 0. 676 0. 689 0. 708 0. 731 0. 762

Units are: Magnetization, cm^g Temperature, K; Magnetic field, Oe.

n
TABLE 20 continued

Sampl* 304-“N Magn«tlxatlon v. <m. T«imp«rat.ur« v. •. Appl i ed F

1

«ld. Unl'ti* “«imj/ gram pag«» 2

H

T 45. 0 50. 0 52. 0 55. 0 60. 0 65. 0 80. 0 114. 0

10 0. 0006 0. 0006 0. 0006 0. 0006 0. 0005 0. 0004 0. 0004 0. 0003

20 0. 0011 0. 0012 0. 0012 0. 0011 0. 0011 0. 0010 0. 0009 0. 0006

30 0. 0017 0. 0018 0. 0019 0. 0017 0. 0017 0. 0016 0. 0013 0. 0010

50 0. 0029 0. 0030 0. 0031 0. 0030 0. 0028 0. 0026 0. 0022 0. 0018

100 0. 0059 0. 0062 0. 0063 0. 0061 0. 0058 0. 0053 0. 0045 0. 0036

200 0. 0119 0. 0125 0. 0127 0. 0123 0. 0115 0. 0108 0. 0092 0. 0074

300 0. 0181 0. 0189 0. 0191 0. 0185 0. 0175 0. 0164 0. 0140 0. 0112

500 0. 0302 0. 0314 0. 0316 0. 0310 0. 0292 0. 0274 0. 0235 0. 0187

667 0. 0400 0. 0417 0. 0426 0. 0412 0. 0389 0. 0366 0. 0311 0. 0251

1000 0. 0613 0. 0640 0. 0644 0. 0630 0. 0596 0. 0560 0. 0478 0. 0382

nil 0. 0660 0. 0713 0. 0715 0. 0699 0. 0664 0. 0621 0. 0533 0. 0426

1250 0. 0767 0. 0798 0. 0807 0. 0787 0. 0748 0. 0700 0. 0595 0. 0478

1429 0. 0071 0. 0909 0. 0917 0. 0895 0. 0848 0. 0797 0. 0661 0. 0546

1667 0. 102 0. 106 0. 107 0. 104 0. 0989 0. 0926 0. 0795 0. 0638

2000 0. 122 0. 127 0. 128 0. 125 0. 119 0. Ill 0. 0952 0. 0763

2500 0. 152 0. 157 0. 159 0. 156 0. 147 0. 139 0. 119 0. 0952

3333 0. 201 0. 211 0. 212 0. 207 0. 197 0. 185 0. 158 0. 127

5000 0. 302 0. 316 0. 316 0. 310 0. 295 0. 276 0. 237 0. 190

7500 0. 455 0. 473 0. 475 0. 464 0. 441 0. 415 0. 356 0. 285

1 0000 0. 609 0. 630 0. 633 0. 619 0. 589 0. 554 0. 475 0. 302

12500 0. 768 0. 787 0. 790 0. 771 0. 735 0. 692 0. 596 0. 478

13000 0. 799 0. 819 0. 821 0. 802 0. 764 0. 721 0. 621 0. 499

239
TABLE 20 continued

Sample* 304~N MagnMt.izat.lon v. «». T^mp^ratur* v. «. Appliod Fi*ld. Uni t«**«mij/gram page* 3

H

T 294. 0

10 0. 0003

20 0. 0005

30 0. 0006

50 0. 0012

100 0. 0025

200 0. 0049

300 0. 0074

500 0. 0124

667 0. 0164

1000 0. 0249

nil 0. 0278

1250 0. 0313

1429 0. 0357

1667 0. 0417

2000 0. 0499

2500 0. 0612

3333 0. 0629

5000 0. 124

7500 0. 167

10000 0. 247

12500 0. 310

13000 0. 323

240
TABLE 21. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AISI 316

Sample 316

V Magnet illation /. s. Temporatur'e v. a. App 1 i od F iold. Unite" emu ' gnam page 1

H " T 3. 0 5. 0 10. 0 15. 0 20. 0 25. 0 30. 0 35. 0

600

^

0. 078 0. 089 0. 089 0. 091 0. 096 0. 103 0. 103 0. 091

667 0. 106 0. 1 14 0. 117 0. 121 0. 128 0. 136 0. 138 0. 121

1000 0. 154 0. 163 0. 173 0. 177 0. 186 0. 202 0. 203 0. 179

nil 0. 170 0. 181 0. 191 0. 196 0. 207 0. 225 0. 227 0. 202

1250 0. 191 0. 201 0. 213 0. 220 0. 232 0. 252 0. 253 0. 226

1429 0. 218 0. 228 0. 242 0. 251 0. 268 0. 287 0. 288 0. 256

1667 0. 255 0. 265 0. 282 0. 294 0. 31

1

0. 335 0. 337 0. 298

2000 0. 303 0- 314 0. 337 0. 347 0. 366 0. 399 0. 399 0. 355

2500 0. 377 0. 391 0. 418 0. 434 0. 456 0. 497 0. 494 0. 437

3333 0. 000 0. 514 0. 554 0. 579 0. 612 0. 664 0. 653 0. 579

5000 0. 000 0. 773 0. 836 0. 877 0. 931 0. 992 0. 956 0. 852

7500 0. 000 1. 175 1. 284 1. 355 1. 423 1. 464 1. 385 1 . 25

1

10000 0. 000 1. 609 1. 778 1. 849 1. 896 1. 898 1. 781 1.617

12500 0. 000 2. 079 2. 289 2. 333 2. 338 2. 297 2. 152 1 . 964

13333 0. 000 2. 238 2. 456 2. 489 2. 478 2. 422 2. 275 2. 076

? -1 Units are: Magnetization, cm g ; Temperature, K; Magnetic field, Oe.

2^1
TABLE 21 continued

Sainp 1 o 316 Mag.not i ir.a L 1 on /. e. T omponatuino v. e. App 1 1 ocl F iolcJ. Unite'-•otnu/gnam pago 2

H T 40. 0 45. 0 50. 0 55. 0 60. 0 65. 0 70. 0 75. 0

'^03 0. 078 0. 069 0. 062 0. 056 0. 051 0. 047 0. 044 0. 040

667 0. 105 0. 093 0. 083 0. 075 0. 069 0. 063 0. 059 0. 054

1000 0. 157 0. 138 0. 124 0. Ill 0. 101 0. 093 0. 087 0. 081

nil 0. 174 0. 154 0. 138 0. 123 0. 113 0. 104 0. 097 0. 090

1?50 0. 195 0. 173 0. 155 0. 139 0. 127 0. 119 0. 110 0. 102

14?9 0. 222 0. 197 0. 176 0. 159 0. 144 0. 133 0. 124 0. 1 16

1667 0. 259 0. 230 0. 205 0. 185 0. 169 0. 155 0. 145 0. 135

2000 0. 309 0. 275 0. 246 0. 221 0. 202 0. 186 0. 173 0. 161

2500 0. 382 0. 343 0. 307 0. 275 0. 253 0. 232 0. 216 0. 201

3333 0. 508 0. 453 0. 404 0. 366 0. 337 0. 309 0. 289 0. 269

5000 0. 757 0. 677 0. 606 0. 546 0. 500 0. 459 0. 429 0. 399

7500 1. 1 17 1. 008 0. 904 0. 817 0. 748 0. 688 0. 645 0. 601

10000 1. 461 1. 325 1. 196 1. 090 0. 997 0. 921 0. 860 0. 803

12500 1. 795 1. 636 1. 491 1. 360 1. 246 1. 150 1. 073 1. 005

1 3333 1 . 907 1. '^34 1. 584 1. 443 1. 327 1. 225 1. 148 1. 073

242
TABLE 21 continued

Samplo 316 Magnet i zcit i on v. e. Toinporaturo v. o. Applied Field. Jn i ts-'-omu/gram page 3

s

' T 80. 0 90. 0 100. 0 110. 0 296. 0

bfl0 0. 038 0. 034 0. 031 0. 029 0. 015

667 0. 051 0. 046 0. 042 0. 038 0. 020

1000 0. 076 0. 068 0. 061 0. 057 0. 030

nil 0. 084 0. 075 0. 068 0. 063 0. 034

1250 0. 095 0. 086 0. 077 0. 071 0. 038

1429 0. 108 0. 097 0. 088 0. 081 0. 043

1667 0. 126 0. 113 0. 103 0. 094 0. 050

2000 0. 151 0. 135 0. 123 0. 113 0. 060

2500 0. 188 0. 169 0. 152 0. 140 0. 075

3333 0. 251 0. 225 0. 202 0- 187 0. 100

5000 0. 374 0. 333 0. 303 0. 281 0. 150

7500 0. 560 0. 503 0. 456 0. 421 0. 225

10000 0. 751 0. 672 0. 609 0. 560 0. 300

12500 0. 937 0. 836 0. 762 0. 702 0. 374

1 3333 0. 999 0. 890 0. 811 0. 748 0. 396

243
TABLE 22. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AISI 316L

'amp 1 e 316 L Magniatlzation v. e. Temperature v. e. Appl 1 ed Field. Unitedemu/gram page 1 H T 4. 2 5. 3 10. 0 15. 0 20. 0 25, 0 30. 0 35. 0

50 0. 0121 0. 0079 0. 0081 0. 0088 0. 0096 0. 0101 0. 0096 0. 0082

100 0. 0198 0. 0150 0. 0160 0. 0177 0. 0188 0. 0202 0. 0199 0. 0170

^00 0. 0364 0. 0310 0. 0320 0. 0341 0. 0367 0. 0405 0. 0399 0. 0343

500 0. 0031 0. 0755 0. 0796 0. 0852 0. 0913 0. 101 0. 100 0. 0862

667 0. 108 0. 100 0. 105 0. 112 0. 120 0. 134 0. 133 0. 1 15

1000 0. 159 0. 152 0. 161 0. 169 0. 183 0. 204 0. 199 0. 173

nil 0. 178 0. 169 0. 178 0. 189 0. 203 0. 226 0. 222 0. 192

1250 0. 198 0. 191 0. 201 0. 213 0. 228 0. 253 0. 249 0. 217

1429 0. 226 0. 217 0. 231 0. 242 0. 261 0. 290 0. 283 0. 246

1667 0. 263 0. 254 0. 269 0. 282 0. 305 0. 336 0. 328 0. 287

2000 0. 316 0. 305 0. 318 0. 340 0. 367 0. 405 0, 394 0. 343

2500 0. 394 0. 385 0. 400 0. 425 0. 460 0. 504 0. 489 0. 427

3333 0. 525 0. 514 0. 539 0. 570‘ 0. 619 0. 672 0. 64 4 0 567

5000 0. 791 0. 777 0, 819 0. 875 0. 949 0. 994 0. 939 0. 836

7500 1. 212 1. 190 1. 270 1. 346 1. 440 1. 447 1. 350 1. 218

1 0000 1. 642 1. 645 1. 767 1. 855 1. 893 1. 855 1. 730 1. 578

12500 2. 131 2. 143 . 2. 269 2. 319 2. 306 2. 231 2. 081 1. 918

13000 2. 219 2. 256 2. 357 2. 407 2. 382 2. 306 2. 156 1. 980

Units are: Magnetization, cm^g”^ ; Temperature, K; Magnetic field, Oe.

244
TABLE 22 continued

Sample 316—L Magnetization v. e. Temperature v. e- Applied Field. Un 1 te^emu/grani page 2. H T 40. 0 50. 0 60. 0 80. 0 181. 0 286. 0

50 0. 0072 0. 0055 0. 0045 0. 0034 0. 0018 0. 0014

100 0. 0147 0. 0115 0. 0094 0. 0070 0. 0039 0. 0029

200 0. 0297 0. 0233 0. 0189 0. 0139 0. 0078 0. 0059

500 0. 0748 0. 0589 0. 0481 0. 0358 0. 0197 0. 0149

667 0. 0990 0. 0781 0. 0637 0. 0474 0. 0256 0. 0198

1000 0. 150 0. 118 0. 0968 0. 0721 0. 0387 0. 0300

nil 0. 167 0. 132 0. 108 0. 0801 0. 0425 0. 0335

1250 0. 188 0. 149 0. 121 0. 0901 0. 0485 0. 0375

1429 0. 214 0. 168 0. 138 0. 103 0. 0552 0. 0430

1667 0. 248 0. 196 0. 161 0. 120 0. 0643 0. 0500

2000 0. 300 0. 236 0. 193 0. 144 0. 0775 0. 0600

2500 0. 374 0. 295 0. 241 0. 179 0. 0969 0. 0751

3333 0. 498 0. 394 0. 322 0. 239 0. 129 0. 100

5000 0. 740 0. 589 0. 484 0. 360 0. 194 0. 151

7500 1. 091 0. 880 0. 725 0. 542 0. 292 0. 226

1 0000 1. 426 1. 164 0. 966 0. 722 0. 390 0. 302

12500 1. 755 1. 445 1. 207 0. 906 0. 489 0. 379

13000 1. 805 1. 502 1. 255 0. 944 0. 510 0. 395

245
TABLE 23. MAGNETIZATION TEMPERATURE (T) AND FIELD (H)

DEPENDENCES OF AUS 330

Sample AWS 330 Magnetization V. e. Ternperature v. e. Appl ied Field. Un i t e"*emu/ gram page

H

T 4. 2 10. 0 15. 0 20. 0 25. 0 30. 0 35. 0 38. 0

10 0. 120 0. 128 0. 138 0. 144 0. 142

20 0. 238 0. 280 0. 318 0. 323 0. 318

30 0. 367 0. 440 0. 488 0. 502 0. 490

50 0. 233 0. 619 0. 619 0. 722 0. 803 0. 814 0. 786

100 0. 588 1. 144 1. 275 1. 412 1. 526 1. 532 1. 492

200 1. 412 2. 186 2. 357 2. 517 2. 665 2. 642 2. 568

300 2. 380 3. 086 3. 268 3. 439 3. 570 3. 490 3. 382

500 1. 435 4. 059 4. 498 4. 634 4. 771 4. 828 4. 669 4. 526

667 2. 010 5. 039 5. 397 5. 494 5. 591 5. 597 5. 392 5. 227

1000 5. 722 6. 644 6. 752 6. 804 6. 815 6. 730 6. 485 6. 286

nil 5. 961 7. 231 7. 140 7. 174 7. 174 7. 003 6. 764

1250 6. 223 7. 572 7. 515 7. 515 7. 515 7. 345 7. 060

1429 6. 530 7. 971 7. 971 7. 971 7. 914 7. 743 7. 458

1667 6. 889 8. 483 8. 483 8. 483 8. 369 8. 199 7. 857

2000 7. 743 9. 110 9. 110 9. 053 8. 939 8. 711 8. 369

2500 8. 825 9. 850 9. 850 9. 736 9. 622 9. 337 9. 053

3333 10. 362 10. 875 10. 818 10. 704 10. 533 10. 248 9. 907

5000 12. 127 12. 298 12. 127 12. 013 11. 785 11. 501 11. 159

7500 13. 664 13. 721 13. 550 13. 380 13. 152 12. 867 12. 526

10000 14. 746 14. 746 14. 518 14. 348 14. 120 13. 892 13. 494

12500 15. 543 15. 543 15. 315 15. 145 14. 917 14. 632 14. 291

13200 15. 771 15. 714 15. 543 15. 372 15. 145 1 4. 860 14. 518

Units are: Magnetization, cm^g"^ ; Temperature, K; Magnetic field, Oe.
TABLE 23 continued

Samplo

H

AWS 330 Magnet i zat i on V. a. T«rriporatur'® v. a. Appl 1 ad Field. Unit«“«mu/gr'

T 40. 0 45. 0 50.0 55. 0 65. 0 187. 0 294. 5

10 0. 136 0. 121 0. 110 0. 0945 0. 0678

20 0. 309 0. 279 0. 240 0. 203 0. 132

30 0- 471 0. 430 0. 366 0. 307 0. 197

50 0. 769 0. 692 0. 585 0. 494 0. 313 0. 0251

100 1. 435 1. 287 1. 105 0. 928 0. 597 0. 0380

200 2. 477 2. 220 1. 913 1. 623 1. 082 0. 0586

300 3. 262 2. 955 2. 562 2. 198 1. 503 0. 0778

500 4. 395 3. 997 3. 519 3. 097 2. 215 0. 113 0. 0427

667 5. 084 4. 657 4. 168 3. 689 2. 710 0. 143 0. 0540

1000 6. 143 5. 705 5. 130 4. 646 3. 570 0. 200 0. 0752

nil 6. 422 5. 978 4. 913 3. 826 0. 219 0. 0826

1250 6. 747 6. 297 5. 215 4. 094 0. 242 0. 0922

1429 7. 105 6. 667 5. 563 4. 412 0. 270 0. 103

1667 7. 515 7. 066 5. 995 4. 788 0. 310 0. 118

2000 8. 028 7. 567 6. 491 5. 301 0. 365 0. 139

2500 8. 654 8. 199 7. 145 5. 944 0. 446 0. 171

3333 9. 565 9. 053 8. 028 6. 832 0. 583 0. 224

5000 10. 818 10. 362 9. 337 8. 199 0. 854 0. 330

7500 12. 184 11. 729 10. 761 9. 622 1. 253 0. 488

1 0000 13. 152 12. 753 11. 785 10. 704 1. 628 0. 647

12500 14. 006 13. 550 12. 639 11. 558 2. 044 0. 803

13200 14. 177 13. 778 12. 867 11. 785 2. 141 0. 848

247
6

rsa ISl cs 3a 3a tsi <39 cs ca CS3 cs 3a <33 <33 cs ca cs CS3 cs <33 <S <39

rvj 03 OO <33 r\j <X3 Applied Field (Oe)

FIGURE 5. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AISI TYPE 310 S
Magnetization

(emu/g)

Applied Field (Oe)

FIGURE 6. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AISI TYPE 304

249

16000
FIGURE 7. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AISI TYPE 304L

2U0
Magnetization

(emu/g)

Applied Field (Oe)

S3 S3 S3 S3 S3 S3 S3 S3 S3

S3 S3 S3 S3 S3 S3 S3 S3

S3 S3 S3 S3 S3 S3 S3 S3

CNJ CO OO ca C%J CO

FIGURE 8. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AISI TYPE 304N

251
csa (S3 (S3 (S3 (S3 (S (S3

csa (S3 (S3 (S3 (Si (S3 (Si (S3 isa S3 (S3 S3 S3 (S3 (S3 (S3 r\j ->3- CO CO (S3 OJ (O

Applied Field (Oe)

FIGURE 9. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AISI TYPE 316

252
Magnetization

(emu/g)

i

Add! led F:;ld (Oe)

I t

FIGURE 10. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AISI TYPE 316L

2o3
Applied Field (Oe)

FIGURE 11. MAGNETIZATION FIELD AND TEMPERATURE DEPENDENCES

OF AWS 330

CS3 csa CS3 csa (SI (33 SJ (S3 csa CSl cs (33 ISi (SI (SJ (33

ISJ C53 (S3 (33 (33 (S3 (S3 OsJ 'CJ* CO CO (33 (NJ CO

2S4
Comments on the Magnetization Results

The interpretations of the significances of the several types

of magnetization signatures encountered in Figures 5 through 11 are given in Table 24.

TABLE 24. MAGNETIZATION SIGNATURES

Characteri Stic Example Interpretation

concave

downwards 310 S superparamagneti sm

1 i near over much of H,T

range

304

316

paramagneti sm over that range

concave

upwards

304

316 at

low temp.

result of the matrixinduced freezing of the

giant moments

extreme downwards

concavity when

unexpected

304L ferromagnetic

contamination

maximum in M(T) 304

316

AWS 330

spin glass

transition

An elaboration of these observations will be presented elsewhere.

Magnetic Susceptibility Temperature Dependences

Derived from Magnetization Studies

Values of static magnetic susceptibility (M/H) were computed at

low (i.e. 500 or 1000 Oe), intermediate, and high (i.e., about 13,000 Oe)

magnetic fields for all of the alloys investigated. The results of this

BATTELLE — COLUIVIBUS

255
. &

work, displayed graphically in Figures 12 through 17 show clearly the

' '’ ®‘*'‘^'6cts of superparamagnetism, spin glass condensation, and these effects'srnmh-i noH

' combined.

i: •it

-fi"'

’s

' •:•

.

m.

m ';>'’ .•SL,;

'

-' f •'.V'iS -u-C: •'isSm

r- .: ^ ' ; ‘ .if' .T:'*fS.L.‘- '.r-* , KJX. j

?

p«- w.:yi

^f5K

«9>v<il

•"V' 5>t*

i’m

r..> “ . •, r"&\ l•l^!=,V^ - ^. »

MU ,j; ''V,

. m'

> 'iR^?S

,., -iss' *
006

0

<D +> CO CO 0 I CD S

CD

»—* (T) c

l4-0

CD Tl 0O

0 0

o o

0) o

0

4^ CO

0

CLX

>x

h- X

(S CS3 Q ro s C3 csa fo

in <3 ea CO ^ cn 00

*

4> »

«>

+1

1*

ca <3 ca 13 C3 — * (3 ca s (3 (3 (3 (3 C3 (3

f\J (3 CD CD OJ

082

092

0t^2

022

002

081

09 T

0^1

02 T

001

08

09

02

0

0

L

D

+>0

L

0

Q_E0

(30-6) /nujeo JO X ui Xr^x •[ xcjx-^cdeosng

257

FIGURE

12.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

AISI

310S

AT

LOW,

INTERMEDIATE,

AND

HIGH

MAGNETIC

FIELDS
100

+ * 03£

<+-0

0 <6 T) CO —

1

9 9

9 •i 0 9 9 3

(0 4- o O O

0 3 3 •—

1

00 +> 3 3 3 3

z 0 3 3 3 3 r-t 0 in 3 3 00

0 Q_• X •—1 in f-H +> X X

CO (— X • + i *

+ 4

14

+• 4

• + I* + i*

• + 4

• + m

• + 4

• + 4

• +

• •% %

(3

O)

3

CO

13 3

to 3in

3 3

CO 3

rj

3 3

082

092

0t’2

022

002

081

0

L

091 3

0

0^1 CLE

9

021 ^

001

08

09

01^

02

0

(00-6) /niueoooTiu X'^ t

j xq t-^deosng

V

FIGURE

13.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

AISI

304

AT

LOW,

INTERMEDIATE,

AND

HIGH

MAGNETIC

FIELDS
t4- —< 0

(9® «

+> T)

CO >-• 2: 0 (9 (9 0 (9 0 1 ‘H O O o O

u» 4-

(9 IS Q IS IS

no +> CS IS cs IS C 0 «—

«

IS IS IS

(9 r-( in no 0.x .—

1

+> X X

CO )— x . + 1 *

« •

-M • Wt* <

»4i IS

(D

4H • mM

-M•MM • -KM

4 -' ISin

—4i IS

mmt, cs

ro

> I 4i IS IS

f\j '

^eg_6) /nuieouoTiu X-:^T

j iqi -:).d0O3n2

me

082

092

0t^2

022

002

081

0

L 091 2

L

0) 0t^T CLE

(9 021

""

001

08

09

0t^

02

0

IS

259

FIGURE

14.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

AISI

304N

AT

LOW,

INTERMEDIATE,

AND

HIGH

MAGNETIC

FIELDS
0

ID LO

COCD (O

U-0

CO

00

0

-p cn

0

o

0

0

Q_X X

ID O

Q

ID 0 O

o

G3 no (3 Q G3 00 ID (3 (3 00

(3 ID m

30E

1382

392

022

002

* *

« « *

+1*

+1 + 1 •+ i •+ I *

•+

1

• +

• -H hi• + .V* 4 . 4 1 1 (— 1 C3 ID 3 ID 3 LD 3 ID 3

!3 X LD 00 3 X ID 00 OJ f

H

08 T

091

0t^T

021

001

08

09

0t^

02

0

0

L

DP

0

L

IDQ_E

ID

(9q_B)

/

niueouoTUJ Xq, t [ r q

t -^ cJeosng

260

FIGURE

15.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

AISI

316

AT

LOW,

INTERMEDIATE,

AND

HIGH

MAGNETIC

FIELDS
IT) GJ LD ca IT) IS tn S5 LO G5 ro SJ IT) CM Q m (M

(M CM « 4 f—< r

(0q_B) /nuieo-jo tuj X-^TTrciT ‘rj.cisosng

30£

382

392

3t^2

022

002

08T

091

0t^I

021

301

08

09

QV

02

0

0L

D

-iJ0L

(5CLEO

2C1

FIGURE

16.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

AISI

31

6L

AT

LOW,

INTERMEDIATE,

AND

HIGH

MAGNETIC

FIELDS
Stainloet?

Steel

Typo

AWS

330

4-0

(D 0 0 COO oo

o o o o O O O C3

•—< C3 O •—* EJ no

+ *

me

mz

092

0V2

022

002

081

091

0t^I

02 T

00T

0

L

D

4>0

L

(PCLE

0

++

+

+

+

+

40 40

40

40

40

40

4=0 *0

*00

*0

08

09

\3V

02

0

o o o o o C3 O O

o o o G5 o O o o

o o o o o O o o

UD C\J o CD CO ru

•—

1

/niueo-jo / TUJ / I TpT :i,deosnc

262

L

FIGURE

17.

MAGNETIC

SUSCEPTIBILITY

TEMPERATURE

DEPENDENCES

OF

AWS

330

AT

LOW,

INTERMEDIATE,

AND

HIGH

MAGNETIC

FIELDS
CONCLUSION

As a result of the magnetization study we have encountered

in four types of austenitic stainless steel, selected for their metallurgical (low temperature strength and stability, and weldability) rather than

magnetic interest four classes of magnetic material; they are:

Superparamagnetic (SPM)

Spin glass

Weakly SPM + spin glass

Strongly SPM + spin glass

AISI 310 S

304

317

AWS 330

Metal physics interpretations of these characteristics will be

advanced in forthcoming research articles.

Magnetic Property Model

As a result of this study the following model for the magnetic

properties of austenitic stainless steels is offered.

1. Stainless steels are not antiferromagnetic in the

long-range ordered sense

2. At elevated temperature they are paramagnetic

3. As the temperature is decreased pre-existing chemical

clusters acquire magnetic order (i.e. giant moments)--

ferromagnetic in the low-Mn steels

4. The matrix within which the clusters are imbedded begins

to undergo spin glass condensation, and finally does so at some temperature at which there is a susceptibility

and magnetization cusp

5. Cluster-matrix interaction tends to induce a frozen random

order into the giant moment system but SPM exists well

below Tg.

BATTELLE"— COLUIVIBUS

263
6. At very low temperatures the magnetization of the

superparamagnetic (cluster) system is anomalously low

due to induced "viscosity". As H increases magnetization of the giant cluster system increases more rapidly than H leading to an upwards concavity in

M(H). This effect is also responsible for the exagerated

shift to low temperatures of the susceptibility maximum

in systems with extreme superparamagnetism.

BATTELLE — COLUIVIBUS

264
LIST OF TABLES

Table

Table

Table

Table

Tabl e

Table

Table

Table

Tabl e

Table

Table

Table

Table

Tabl e

Tabl e

Table

Tabl e

Table

Tabl e

Table

1. Temperature Ranges of the Susceptibility and

Magnetization Studies

2. Vendors and Available Specifications of the

Alloys Investigated

3. Vendor-Supplied Composition of Alloys Listed in

Wgt Percent

4. Degree of Etching and Specimen Weights in Magnetic

Susceptibility and Magnetization Measurement . . .

5. Magnetic Susceptibility of AWS 330. Graphically

Extrapolated x- Values for the Temperature Range

414. 5-129. 5K .“

6. Magnetic Susceptibility of AISI Type 304 in the

Temperature Range 8.7-65.6K

7. Index of Computer-Produced Susceptibility Temperature

Dependence Data Tables

8. Magnetic Susceptibility Temperature Dependence of

AISI 310S Over the Temperature Range 77-414K . . .

9. Magnetic Susceptibility Temperature Dependence of

AISI 304 Over the Temperature Range 77-408K . . . 10. Magnetic Susceptibility Temperature Dependence of

AISI 304L Over the Temperature Range 77-410K . . . 11. Magnetic Susceptibility Temperature Dependence of

AISI 304N Over the Temperature Range 77-403K . . . 12. Magnetic Susceptibility Temperature Dependence of

AISI 316 Over the Temperature Range 77-407K . . . 13. Magnetic Susceptibility Temperature Dependence of

AISI 316L Over the Temperature Range 77-413K . . .

14(a). Magnetic Susceptibility Temperature Dependence of

Inconel 625 Over the Temperature Range 9-80K . . .

14(b). Magnetic Susceptibility Temperature Dependence of

Inconel 625 Over the Temperature Range 77-41 2K . .

15. Coefficients of the Susceptibility Temperature

Dependences of Some 300-Series Stainless Steels and

Inconel 625

16. Index of the Temperature and Magnetic Field

Ranges Covered

17. Magnetization Temperature (T) and Field (H)

Dependences of AISI Type 31 OS

18. Magnetization Temperature (T) and Field (H)

Dependences of AISI 304

19. Magnetization Temperature (T) and Field (H)

Dependences of AISI 304L

265
LIST OF TABLES

(Continued)

Table 20. Magnetization

Dependences

Table 21. Magnetization

Dependences

Table 22. Magnetizati on Dependences

Table 23. Magnetization

Dependences

Table 24. Magnetization

Temperature (T) and Field (H)

of AISI 304N

Temperature (T) and Field (H)

of AISI 316

Temperature (T) and Field (H)

of AISI 316L

Temperature (T) and Field (H)

of AWS 330

Signatures

LIST OF FIGURES

Figure 1 (a). Magnetic Susceptibility of AWS 330. Reciprocal Field

Plots for the Temperature Range 414. 5-196. 5K. The

Gradual Development of "Ferromagnetic Slope" is

Evident

Figure 1 (b). Magnetic Susceptibility of AWS 330--Continued.

Reciprocal Field Plots for the Temperature Range

185-77K. The onset of Superparamagnetism at

Temperatures Near lOOK is Evident

Figure 2. Magnetic Susceptibility Temperature Dependence of

AISI Type 304, at Two Magnetic Fields Levels,

Within the Temperature Range 8.7-65.6K

Figure 3. Algorithm for the Computer Reduction of Susceptibility

Force Balance Data

Figure 4. Magnetic Susceptibility Temperature Dependences

for AISI 310 S, AISI 304, AISI 316, AWS 330 and

In 625

Figure 5. Magnetization Field and Temperature Dependences

of AISI Type 310 S

Figure 6. Magnetization Field and Temperature Dependences

of AISI Type 304

Figure 7. Magnetization Field and Temperature Dependences

of AISI Type 304L

Figure 8. Magnetization Field and Temperature Dependences

of AISI Type 304N

Figure 9. Magnetization Field and Temperature Dependences

of AISI Type 316

2GC
LIST OF FIGURES

(Conti nued)

Figure 10

Figure 11

Figure 12

Figure 13

Figure 14

Figure 15

Figure 16

Figure 17

. Magnetization Field and Temperature Dependences

of AISI Type 316L

. Magnetization Field and Temperature Dependences

of AWS 330

. Magnetic Susceptibility Temperature Dependences of

AISI 310S at Low, Intermediate, and High Magnetic

Fields

. Magnetic Susceptibility Temperature Dependences of

AISI 304 at Low, Intermediate, and High Magnetic

Fields

. Magnetic Susceptibility Temperature Dependences of

AISI 304N at Low, Intermediate, and High Magnetic

Fields

. Magnetic Susceptibility Temperature Dependences of

AISI 316 at Low, Intermediate, and High Magnetic

Fields

. Magnetic Susceptibility Temperature Dependences of

AISI 31 6L at Low, Intermediate, and High Magnetic

Fields

. Magnetic Susceptibility Temperature Dependences of AWS

330 at Low, Intermediate, and High Magnetic Fields . . .

267
APPENDIXES A AND B

Discussions of the influence of a' martensite on the magnetic susceptibility and magnetization of

Type 304 alloys.

APPENDIX A MAGNETIC EVALUATION OF MACHINING INDUCED

DEFORMATION TRANSFORMATION IN AUSTENITIC

STAINLESS STEELS

APPENDIX B MAGNETIZATION STUDIES OF MARTENSITE-BEARING

AISI TYPE 304 STAINLESS STEELS

BATTELLE — COLUMBUS

268
APPENDIX A

MAGNETIC EVALUATION OF MACHINING INDUCED DEFORMATION

TRANSFORMATION IN AUSTENITIC STAINLESS STEELS

E. W. Col lings

Accepted for publication in CRYOGENICS

ABSTRACT

Austenitic stainless steel alloys, particularly those with high

Mj temperatures are prone to transform to ferromagnetic martensite

during deformation. Prior to studying their magnetic properties, therefore, care must be taken to limit sample-preparation-induced damage

to the surface layers, and then to remove these by severe etching before

making the measurements. A magnetic technique for evaluating deformationinduced transformation is described and some experimental results are discussed.

INTRODUCTION

Before engaging on a program to study the low-temperature

magnetic susceptibilities of a series of commercial stainless steel

alloys it was recognized that some members of the series, particularly

those with high (near room- temperature) M^ temperatures such as the

304's [1], would be prone to transform martensi tical ly as a result of

handling and machining during sample preparation. Two courses are open to remedy this situation -- the sample could be heat-treated in

which case it would not longer be a true representative of the parent

alloy; or it could be etched, a remedy effective only for surface

damage, and one which possesses the possible disadvantage that

BATTELLE — COLUMBUS

269
preferential etching might leave the surface rich in one or more of

the ferromagnetic constituents (Fe or Ni).

Using 304N as sample, we applied the magnetic susceptibility

technique itself to estimating semiquantitatively the degree of transformation which accompanied (a) bulk deformation (b) surface damage due

to cutting, and (c) to examine the effectiveness of controlled etching

in removing surface damage without deleterious side effects.

EXPERIMENTAL

The magnetic susceptibilities of small cube-shaped samples were measured by the Curie (magnetic buoyancy) technique in fields (H)

of between 3.75 and 10.0 kOe (i.e., between 3.0 and 8.0 lO^Am""').

Assuming the sample to contain a mass-fraction f of ferromagnetic

contaminent of specific magnetization a, the magnetic susceptibility

X, is given by

X = (l-f)^T + f a/H .

Thus, provided that the ferromagnetic component is always saturated

in the measuring field (H ^3.75 kOe, i.e., 3.0.10^Am such that

a = const., a plot of x versus H

^ is linear with intercept (1-f)^^,

where -j is the susceptibility of the matrix and slope f a. Thus, knowing

a one obtains f the fraction of ferromagnetic material present. In this work we assume for a the saturation moment of a-Fe, viz. 221.9

3 -1 Oe.cm g ,* thus acquiring a lower limit for the amount of transformed

martensite. In the present measurements the standard error in slope

-2 3 -1 IS +0.14.10 Oe.cm .g .* "Zero Slope" thus represents an a-Fe content

of +6.3 ppm.

The samples, which were removed from supplied stock material

using a liquid-sprayed carborundum cut-off wheel, and which were

* 1 Oe.cm^.g""' = (1/4 tt) A.m^.kg"\

BATTELLE — COLUMBUS

270
measured both before and after etching with "glyceregia" ,** can be

defined according to the following scheme:

Sample 1, (a) 24.75 mg, heavily deformed, measured; (b) etched

to 18.03 mg, remeasured.

Sample 2, (a) 20.70 mg, as-cut, measured; (b) etched to 13.25 mg,

remeasured.

Sample 3, (a) 5.70 mg, as-cut, measured; (b) etched to 4.37 mg,

remeasured.

Thus, in Sample 1, we were dealing with both bulk and surface

deformation, while in Samples 2 and 3 we examined surface effects

only, at two different surface/volume ratios.

RESULTS

The results, in the format x versus H"^ are presented in

Figures 1(a) and (b) and summarized in Table 1,

DISCUSSION AND SUMMARY

Stainless steels such as the 304' s with near-room-temperature

M^'s [1] are subject to martensitic transformation during sample

preparation. Local bulk plastic deformation in holding grips, and

surface deformation in cutting or grinding induce sufficient transformation

to cause serious measurement errors unless adequate precautions are

taken, flaterial heavily deformed by compression can yield some 0.1 wgt%

of transformation product, sufficient to increase the susceptibility

at 5 kOe (4.10 Am ) by a factor of 7 above the true bulk value

(from 26.5 to 182 10 ^ cm^ g

^ (10~^m^kg at room temperature)

,

Figure 1(a). With small (*“ 6 mg) wheel -cut samples of 304N the a-Fe-

** 10 ml HNO^, 20-50 ml HCL, 30 ml glycerol

BATTELLE — COLUIVIBUS

271
5 -1 equivalent contamination is 0.03 wgt%, which raises the 5 kOe (4.10 Am ) susceptibility by a factor of 1.5, Figure 1(b). Naturally, the cuttinginduced relative surface damage decreases as the surface/volume ratio

decreases.

Measurements made at a single value of magnetic field cannot,

therefore, be relied upon to yield even an approximate value of

magnetic susceptibility. Field-dependent measurements must be undertaken

in order to determine, by extrapolation to H

^ = 0, the true bulk

value, and to monitor the progress of any surface-damage-removal procedure.

For some types of austenitic stainless steels, e.g., 3105, the magnetic

susceptibility, as a result of superparamagneti sm [2], becomes more

^nd more field-dependent as the temperature is lowered, others transform

structurally or magnetically. Thus, in order to be able to examine

these phenomena properly all sources of spurious field-dependence at room temperature must be eliminated prior to commencing a series of lowtemperature experiments.

ACKOULEDGEMENTS

Sample preparation and magnetic measurements were carried out

by R. D. Smith. The research was funded by the U.S. Department of

Energy through the National Bureau of Standards, Boulder, CO, Dr.

F. R. Fickett, Program Manager.

REFERENCES

[1] L.A.A. Warnes and H.W. King, Cryogenics, 16^, 659 (1976)

[2] E.W. Collings, F. J. Jelinek, J. C. Ho and M.P. Mathur

Advances in Cryogenic Eng., 159 (1977).

BATTELLE — COLUIV1BUS

272
Q

UJO

q;oQ

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h-c

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LU

ooc

LU_l

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I—

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DCct

-M CM CM C\J cd- CM

4->

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oo LO C\J CO

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+ 1 + 1 a 1— '—

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03 OO I cr

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«=C

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273

Oe.cm'^.g =

(1/4it )
FIGURE A

FIGURE A (a) Reciprocal-field plots of magnetic susceptibility

for heavily deformed 304N (greater than 50% by compression), and

(b) reciprocal -field plots for as-cut and as-cut-plus-etched samples.

Mass reduction of 20-40% by etching resulted in an undetectable slope,

corresponding to a deformation-martensite a-Fe equivalent of

+ 6.3 wgt.ppm. 20% mass reduction of a cube-shaped 6mg sample

_2

corresponds to the removal of surface to a depth of ^'1.5 X 10 mm.

274
Reciprocal Magnetic Field, H MA ' m

0 I 2 3 4 5 6 7 8 9 10 II 12 13

Reciprocal Magnetic Field, H ', MA ' m

0 1 2 3 4 5
APPENDIX B

MAGNETIZATION STUDIES OF MARTENSITE-CONTAINING

AISI TYPE 304 STAINLESS STEEL

In Figures 6 and 8 we note the linearity in M versus H for

304 and 304N. On the other hand. Type 304L seems completely out of

character, with extreme curvature to M(H). This apparent anomaly is

resolved when we note that the high field slopes of M versus H for 304L are similar to the slopes (at corresponding temperatures ) for the other

alloys over the entire field range. Clearly Type 304L contains a saturable ferromagnetic contaminant (c.f. Appendix A).

It was initially supposed that the ferromagnetism noted was a result of cutting-induced surface damage. Accordingly, four more samples were prepared, very heavily etched, and measured at room temperature. As

the following four curves (Figures B1 through B4) indicate, the ferromagnetic component (presumably martensite) was not confined to the surface,

but was distributed throughout the bulk of the specimen as supplied.

Further investigation would involve annealing and so on, but the effect

is in principle solved.

This observation emphasizes that extreme care must be taken

when dealing with marginally stable austenitic stainless steels (i.e.

high M , temperatures), to ensure that the properties being measured

S 5 Q are intrinsic to the alloy and are not being perturbed by the effects

of parasitic precipitates and transformation products.
Magnetization

(eitiu/g)

Applied Field (Oe)

FIGURE B-1 . FIELD DEPENDENCE OF ROOM TEMPERATURE MAGNETIZATION

OF AISI TYPE 304L SAMPLE 1

277

00091
FIGURE B-2. FIELD DEPENDENCE OF ROOM TEMPERATURE MAGNETIZATION

OF AISI TYPE 304L SAMPLE 2

27L
Magnetization

(emu/g)

FIGURE B-3. FIELD DEPENDENCE OF ROOM TEMPERATURE MAGNETIZATION

OF AISI TYPE 304L SAMPLE 3

279
Applied Field (Oe)

FIGURE B-4. FIELD DEPENDENCE OF ROOM TEMPERATURE MAGNETIZATION

OF AISI TYPE 304L SAMPLE 4

280
SOME MAGNETIC PROPERTIES OF THE

"NONMAGNETIC" STAINLESS STEELS

National -Bureau of Standards

2G1

SOME MAGNETIC PROPERTIES OF THE "NONMAGNETIC" STAINLESS STEELS

F. R. Fickett

Thermophysical Properties Division

National Bureau of Standards

Boulder, CO 80303

ABSTRACT

The stainless steels proposed for low temperature service exhibit magnetic

properties which are relatively complex functions of temperature. Welds in

these metals also show complicated behavior. This paper discusses the magnetic

properties of the "nonmagnetic" stainless steels as presented in the literature.

The problems with using the permeability as a design parameter are outlined.

Data are presented on the very low field susceptibility and the remanent magnetization of a large number of stainless steels measured at room temperature and at

4 K with a SQUID (Superconducting Quantum Interference Device) Magnetometer System.

INTRODUCTION

The austenitic stainless steels are, for practical purposes, paramagnetic

materials at room temperature and above. However, as the temperature is lowered

they all undergo transitions to other, more complex, magnetic states. The exact

nature of these low temperature states is still open to some debate, but at present it appears that they form spin glass systems as discussed by Col lings elsewhere in this report. The prior literature indicates that the state might also

be one of anti ferromagnetism or superparamagnetism. The actual state depends

more on the interpretation of the data than it does on the material. The two

measurements used to investigate these magnetic state properties are the variation of the magnetic susceptibility with temperature and the variation of the

magnetization with applied field at fixed temperature. Graphs of these properties for representative austenitic steels are shown in Fig. 1 and Fig. 2.

283
Note that the magnetization at room temperature is linear with field giving

a constant susceptibility (k= M/H)* indicative of true paramagnetism. At 4 K,

however, k is strongly field dependent.

The location of the peak of the susceptibility curve (Fig. 1) depends

strongly on the specific steel and may vary from 30 K to nearly room temperature.

In the measurements shown (Ledbetter and Collings, 1978), any truly ferromagnetic

components have been saturated out so that the data are for the austenitic phase

alone.

While all this is very interesting from a metal physics viewpoint, in engineering applications one must be able to achieve a specific value of the relative magnetic permeability, p^. by an appropriate choice of material for service at

(usually) a fixed temperature. The relative permeability, p^, is simply related

to the susceptibility (in SI units) by

p/ =y/vio = 1 + 0)

Unfortunately , it is conventional to give susceptibility values in units of

emu/g. We call this quantity Xp» the cgs mass susceptibility. Xp must be

3 multiplied by the material density to convert it to x in units of emu/cm

which is, believe it or not, dimensionless. This quantity, multiplied by 4tt

is K. The situation is only slightly less confusing if SI units are used

throughout.

The use of permeability as a specification on stainless steel at low temperatures should only be done with a clear understanding of the implications of

the M-H behavior shown in Fig. 2. The classical definition of p is directly

related to the slope of the line from the origin to the curve and thus can have

*K is the symbol chosen to represent the SI susceptibi 1 i ty, which is dimensionless. The cgs susceptibility, represented by x is also dimensionless, but

K = 4ttx.

284
many values as shown. It is not a meaningful measure in general. However, the

initial value of p, i.e. that measured in fields < 8 kA/m (100 Oe) is a useful

number, i_f the shape of the full curve is kept in mind. Note that this is only

a problem at low temperatures - at room temperature y is a constant with field.

Another problem with the austenitic stainless steels, which is usually

minor, is that truly ferromagnetic regions may exist in the steel due to a

variety of causes, such as surface oxidation, formation of stress-induced or

thermally-induced martensite, and ferrite formation which is particularly a problem in weldments. All of these cause the steel to show a magnetic remanence

after magnetization which (rarely) may be significant in applications.

In this report we describe measurements of the magnetic remanence at room

temperature and the susceptibility at both room temperature and 4 K for a

variety of austenitic steels and some preliminary room temperature results on

weld metals. Unless otherwise indicated, the measurements are made with a very

sensitive magnetometer using a SQUID (Superconducting Quantum Interference

Device) detector. This device is described by Pickett and Sullivan (1974).

All susceptibility measurements are true initial susceptibilities. The base

metal specimens are 20 mm long rods ^ 3 mm square with rounded edges. The weld

metal specimens are 'v 5 mm diameter cylinders, also 20 mm long.

RESULTS

Remanent Moment

The results of room temperature measurements of the remanent magnetization

after application and removal of a 358 kA/m (4.5 kOe) field are shown in Fig. 3.

The as-machined (M) specimens show a significant variation from the annealed

(A) ones, but with no systematic trend. The measurements which fall below the

dashed line are probably of no significance to applications of the steels.

They would give trouble in only the most demanding situations. The high value
for the machined 304 is most likely the result of martensite formed by

deformation.

Susceptibility

The results of measurements of the initial susceptibility of the base

metals at room temperature and at 4 K are shown in Fig. 4. Because this is

a true initial susceptibility measured on previously unmagnetized material,

the ferromagnetic components seen in Fig. 3 will also show up here as an

increase in the susceptibility. They should not be very temperature dependent,

however. It is clear that cooling to 4 K changes the susceptibility, usually

significantly. This is what we would expect from the general behavior shown

in Fig. 1 and, in fact, there is good agreement at both temperatures for the

few cases in the literature where the x vs T curves have been taken to 4 K.

Also, it appears that the machining and annealing operations do not greatly

alter the susceptibility except where ferromagnetic inclusions are apparently

created.

Weld Metal Properties

Ferrite in weld metal is a ferromagnetic material and, as such, will

show up both in the initial susceptibility and remanence measurements. We would

not expect a strong temperature dependence. A preliminary measurement of the

room temperature initial susceptibility of 316 and 31 6L weld metal specimens is

shown in Fig. 5. The percent ferrite is determined by a Magnegages which measures the force required to pull a calibrated magnet from the metal. The susceptibility was determined by a two-coil induction device (Cristie and Symons,

1969) which applies a maximum field of 40 A/m (0.5 Oe). The specimens were

measured transverse to the field and various calibration and correction factors

were not determined in these preliminary measurements and, thus, only relative

values of susceptibility are presented. Initial estimates indicate that,

around 10% ferrite, the relative permeability may be as high as = 2 .
CONCLUSIONS

Austenitic stainless steels are not classical ferromagnets at any level

and there is a danger of creating confusion by using a terminology appropriate

to the ferromagnets. The most meaningful data for completely defining the magnetic properties of these metals at any temperature is a curve of magnetization

versus applied field.

A remanent ferromagnetism due to inclusions and surface damage is found to

exist in many of the stainless steels. The effect is usually negligible, but

should be determined in some of the steels after machining or heat treatment if

the application is at all critical magnetically.

The initial susceptibility is temperature dependent and depends also on the

specific steel. It is not strongly affected by annealing or machining unless

ferromagnetic particle formation occurs.

Weld metals are significantly more ferromagnetic than any of the base

metals due to ferrite formation. More study is needed, particularly to determine

if the magnetic properties are truly temperatue independent as we have assumed.

237
REFERENCES

Christie, K. W. and Symons, D. T. A. (1969). Apparatus for measuring magnetic

susceptibility and its anisotropy. Geological Survey of Canada, paper 69-41.

Efferson, K. R. and Leonard, W. J. (1967). Magnetic properties of some structural materials used in cryogenic applications. Oak Ridge National Laboratory

Report ORNL-4150, pp. 126-132.

Fickett, F. R. and Sullivan, D. B. (1974). Magnetic studies of oxidized impurities in pure copper using a SQUID system. J. Rhys. F: Metal Rhys. 900-905.

Ledbetter, H. M. and Collings, E. W. (1978). Low-temperature magnetically

induced elastic-constant anomalies in three manganese stainless steels. In:

Metal Rhysics of Stainless Steels (TMS-AIME, New York) to be published.

283
( 6 /nuj 0 01) Ainiaiid33sns ssvw

9

-

(6>|/ lu 01) Ainiaild3DSnS ssvw

c o

209

Figure

1.

The

variation

of

magnetic

susceptibility

asa

function

of

temperature

for

the

stainless

steel

Fe-21

Cr-6Mi

-9Mn

.

The

behavior

is

typical

of

nearly

all

austenitic

stainless

steels,

although

the

temperature

of

the

peak

varies

with

the

steel.

From

Ledbetter

and

Col

lings

(1978).
APPLIED

FIELD,

H

(kOe)

(ui/V>l) W ‘N0I1VZI13N9VW

290

Figure

2.

The

magnetization

of

AISI

310

asa

Function

of

applied

field.

The

behavior

shown

is

typical

of

all

austenitic

stainless

steels.

The

dashed

lines

indicate

how

the

relat'^ve

permeability

is

determined.

Data

from

Efferson

and

Leonard

(IQb-?).
ROOM

TEMPERATURE

(6/niua) 1N3W0W lN3NVW3il

oo

CM

I (T)I oI o

un

I o

<X) CO

CM

oLO

O

oo

ro

CO

l-H CO

o

CO

oCO

(O "O O 0)

•I- c

Q.

035^

JOo(Uo : 3

x>

• c

CO ,

0) I <v

-t-J CO

CO

1/1 a;c

03

4-> COu

c

O). -M

CO3

03 M-O

+-> C

O)o

E

O)c

03E

01

s01

ro

O)

s-3

C71

(6>|/^iuV) 1N3W0W 1N3N\/W3a

291

tion

of

a

358

kA/m

(4.5

kOe)

field.

Specimens

labeled

M

are

mach

from

as-received

stock.

Those

labeled

A

have

been

annealed

after

machining.
(5/nuj0) AiniaiidBDSnS SSVW 1VI1INI

(6>|/^U]) A111iaild33sns SSVW IVIIINI

292

Figure

4.

The

initial

susceptibility

at4

K

and

at

room

temperature

for

austenitic

stainless

steels.
Aini811d33sns 3AIlV13a

293

measured

at

room

temperature.

Stainless Steel Welding

295
•i

I

t

'i

jt -A

1

^ I 1

i

f

!
EVALUATION OF STAINLESS STEEL WELD METALS

AT CRYOGENIC TEMPERATURES

National Bureau of Standards

297

Evaluation of Stainless Steel Weld Metals

at Cryogenic Temperatures

H. I. McHenry, D. T. Read and P, A. Steinmeyer

National Bureau of Standards

Boulder, Colorado 80303

The program to evaluate stainless steel weld metals for structural applications in superconducting magnet systems begun in FY 77 and was continued in

FY 78. The objectives of the program are;

1. To investigate the metallurgical factors which affect the toughness

of stainless steel weld metals.

2. To develop a high strength, high fracture toughness filler wire

for welding nitrogen-strengthened stainless steels such as the

21Cr-6Ni-9Mn alloy.

3. To evaluate the fracture toughness of stainless steel welds at

4 K.

During FY 77, the first year of the program, three filler metal companies

participated in the program(l). Arcos Corporation prepared a series of 10

test weldments to study the effect of ferrite content, ferrite morphology and

sensitization on 316L weld metal deposited by the shielded metal arc (SMA)

process. Armco evaluated the influence of nitrogen content on the strength and

toughness of three experimental weld metals at 76 K: 1 3Cr-20Ni-9Mn-l .8Cb

,

20Cr-l 0Ni-9Mn and 18Cr-16Ni-9Mn. Tel edyne-McKay prepared SMA weldments with a

18Cr-34Ni-4.5Mn weld metal, a modified AWS 330 alloy. During FY 78, the Arcos

and Teledyne-McKay weldments were received and the metallurgical and fracture

toughness evaluations were completed. In addition, Armco completed the remaining

work on the experimental alloys and work was started on a filler metal selection

program in cooperation with General Electric's Large Coil Program. This report

summarizes the results obtained in FY 78 and briefly outlines the plans for FY 79.

299
Metallurgical Effects on Toughness

The 31 6L welds supplied by Arcos were used to study the influence of ferrite

content, ferrite morphology and sensitization on toughness at cryogenic temperatures,

particularly 4 K. The metallurgical evaluation was conducted by the Colorado

School of Mines and the results are presented in the attached report by Elmer,

Steinmeyer and Olson. The fracture toughness tests were conducted by NBS and the

results are presented in the attached report by Read, McHenry, Steinmeyer, and

Thomas. The results clearly indicate that ferrite and sensitization cause

significant reductions in toughness at 4 K.

High Strength Filler Wire Development

In the first year report(l), Armco presented the results of a program to

evaluate the influence of nitrogen on the strength and toughness of three

experimental filler wires at 77 K. The results, summarized in Figure 1, indicate

that the strength of each alloy increases and the toughness decreases with increasing

nitrogen content. Each of the alloys had marginal toughness at 77 K even at the

lowest nitrogen level, and conseguently plans for further testing at 4 K were

cancelled.

The welds evaluated in the Armco Program were deposited by the SMA process.

Stainless steel welds deposited by the SMA process generally have less toughness

than similar welds deposited by either the gas tungsten arc (GTA) or gas metal

arc (GMA) processes. For comparative purposes, a pulse-mode GMA weld was prepared

for each experimental alloy. Nitrogen was kept to a minimum; however, the

nitrogen levels were relatively high due to the nitrogen content of the wire.

The chemistries of the as-deposited weld metals are shown in Table 1 and the

mechanical properties at room temperature and 77 K are summarized in Table 2.

The 1 3Cr-20Ni-9Mn-2Cb and 20Cr-10Ni-9Mn alloys had marginal toughness at 77 K

and thus, further tests at 4 K are not planned. The 18Cr-l 6Ni -9Mn alloy had a

300
promising combination of strength and toughness at 77 K and consequently NBS

will conduct fracture toughness tests at 4 K during FY 79.

Evaluation of AWS 330

The AWS 330 test plates prepared by Tel edyne-McKay and described in last

year's report(l) were fracture toughness tested at room temperature, 76 K and

4 K. The results, summarized in Table 3, indicate that AWS 330 has relatively

good strength and toughness at 4 K. Microfissures were observed on the fracture surface; however, the microfi'ssuring tendencies were not characterized or

studied.

Electrode Selection Tests - 308L vs 31 6L

A series of test weldments were prepared by General Electric using the

SMA process for an evaluation of three candidate electrodes: E308L, E316L

(N = .044%) and E316L (N = .071%). The base metal was 316LN, one-inch-thick

plate. Specimens were sent to NBS where the following tests were conducted on

each weld: tensile properties at room temperature and 4 K and fracture toughness tests at 4 K. In addition GE conducted Charpy impact tests on each weld

metal at room temperature and 77 K.

The welding procedures, and the weld deposit chemistries are summarized

in Tables 4 and 5. Tensile tests were conducted on al 1 -wel d-metal specimens

(specimen axis parallel to the weld) and on cross-weld specimens which included base metal, heat affected zone and weld metal in the test section. The

results are summarized in Table 6. All of the cross-weld specimens failed in

the weld indicating that the weld had lower strength than either the 316LN

base plate or the heat affected zone. However, at 4 K, the strength level of

the 316L weld metal was comparable to that of the 316LN base metal, particularly at the .071 N level where the joint efficiency was 89%.

301
The fracture toughness data at 4 K, presented in Table 7, indicate that

308L has better 4 K fracture toughness than 31 6L. The data of Table 7 may be

compared with the results for the Arcos weldments, in the attached report by

Read, McHenry, Steinmeyer, and Thomas. The toughness of 31 6L weld metal at

4 K measured on the Arcos weldments was 180 MPav^ (164 ksi/Tn at the zero

ferrite level. In contrast, the GE 316L welds had values of 118 MPat/m

(107 ksi/Tn) (N = .044%) and 95 MPa/m (86 ksi/Tn) (N = .071%). The only

identified difference in the Arcos and GE welds other than minor variations

in chemistry and weld procedures, was the electrode coating; titania coatings

were used by Arcos and lime coatings were used by GE. Further work is needed

to understand the cause of the large differences in toughness.

Further confusion arises when the 77 K Charpy impact data obtained by GE,

Table 8, are compared with the data on the Arcos welds. The GE 316L welds had

significantly higher Charpy impact values at 77 K (51J (38 ft-lb) at .044%N

and 45J (33 ft-lb) at .071%N) than the Arcos welds (35J (26 ft-lb) at .04%N).

In comparing the GE welds, 308L had the best toughness at 4 K and lowest Charpy

impact energy (31 J (23 ft-lb) at 77 K. Thus, use of a Charpy test at 77 K to

estimate fracture toughness at 4 K does not appear useful

.

Future Work

A. Multi-process study of 316L. Test weldments (AWS 316L filler metal,

AISI 304L base plate) will be prepared by a subcontractor using each of the

following processes: gas tungsten arc, gas metal arc and submerged arc.

Initially, the welds will be evaluated by fissure-bend tests, room temperature tensile tests and Charpy tests at 76 K. Subsequently the welds will be

evaluated at 4 K using tensile and fracture toughness test methods. All

welding and initial testing will be done in FY 79. Testing will be completed

in FY 80.

302
C. Aluminum Welding Program. Several fabricators will be asked to prepare

test welds in 5083-0 aluminum using 5183 filler wire obtained from a source of

the fabricators choice. The weld procedures used and the resultant test welds

will be qualified in accordance with ASME Boiler and Pressure Vessel Code

requirements (Sections VIII and IX). The results will indicate if property

variations at 4 K are significant enough to warrant more stringent qualification

procedures for welds used in liquid helium applications.

D. Large Coil Project Support. Welds supplied by General Electric and

General Dynamics will be evaluated at 4 K to develop data for use in designing

superconducting magnets for the Large Coil Project.

303
Table

1.

Chemical

compositions

of

the

gas

metal

arc

weld

deposits

of

three

experimental

alloys

Data

provided

by

Armco

(weight

percent) !

I'

304
Table

2.

Mechanical

properties

of

gas

metal

arc

weld

deposits

of

three

experimental

alloys

Data

provided

by

Armco

30G
Table

3.

Fracture

toughness

of

AWS

330

weld

metal

shielded

metal

arc

welds

in

one-inch-thick

304L

plate

<u

s-3

-M

03

SO)

. c

<JD CM CM

1— CM 00 0 CM UO

(/) CM r*o CO CO 1— 1—

u—t

CT\ <0 CD 0 LT) 0 03 ro 10 CO CO

Q_ CM ro CO 1— j—

s:

c 0 0 0 0 •1— LT> 0 0 0 0 0 CO f— f— 0 JD 1— 00 UO CO UO 00

f—

O—1

uo ^ E CT3 uo CO CO . 0 0

00 CJ3 00 1—

1

2: CM uo 00 uo 00 r~^ -0 c

UJ 0 "—^ *P“

+-> C

II 03 0

s- •

1—

(J 3 +J

' CD 03 in ^ -1- SCTi CD CD CD 4- 3

CM • • C CD 0 0 •I—

•—I U 4- O C

4-> 0

E c

0 •!-

Jc s- 0 +-> ¥ 4- -r-> c

¥ •I—

.—

»

N —'' »—

-D 01 0 -0 T3 -a T3 03 03 ^¥ 03 03 03 03 »-> > —

s

--—^ +-> +J +-> -l-> 03 1 -0 T3 3 3 3 3 — 03 • O) 0) O) f-^ F 1— f— 3 1— uo

+-> +-> *r" * r~ •t— r“ U -Q

Q. 3 3 "O -a -c -0 1— 3 CD E 1— c c. c 03 0 3 03 *r~ •r~ 3 3 3 3 C_) Q <C

00 -0 -a ''—

^ ¥ ¥ ¥

s.—

^ ¥ ¥

CM CM CO m ¥

CM CO 1 1 1 1

1 1 CQ ec CQ c CO IS 3 3 3

306
Table

4.

Chemical

composition

of

shielded

metal

arc

weld

deposits

in

one-inch-thick

316LN

plate

Data

provided

by

the

General

Electric

Company

(weight

percent)

I

to to

• • • o CO CO r— r—

s- o

to LO tn

• • •

to

C\J o CM OJ ro CO

• • • o o o

o<Ti

'd- Oocri

CM

CM

LO

CM O O CM

• • • o o o

o

LO LO LO

1 '“I'

'

1 _J _i

1 _J CO to to o t 1—

ro CO CO

UJ LU IjU

O)CD

fO

C71

I fOc

cn

<o

307
Table 6. Tensile properties of shielded metal arc weld deposits

in one-inch-thick 316LN plate

Specimens supplied by the General Electric Company

Yield Ultimate Elon- Reduction

Temperature Strength Strength qation of Area

Material Specimen K MPa (ksi; MPa (ksi) ( in 1")^^ %

316LN Base Plate 295 279 (40.5) 615 (89.3) 57 72

Base Plate 4 1025 (149) 1388 (201) 35 50

E308L-15 All-Weld 295 354 534 19 72

All-Weld 295 408 543 20 12

All-Weld Avg 381 (55.3) 538 (78.1) 20 67

All -Weld 4 598 1358 45 44

All -Weld 4 709 1357 38 50

All-Weld Avg 654 (94.9) 1358 (197) 42 47

Cross-Weld 295 407 (59.1) 568 (82.4) 27 57

Cross-Weld 4 691 1376 31 43

Cross-Wel

d

4 699 1401 30 49

Cross-Weld Avg 695 (101) 1388 (201) 31 46

E316L-15 All -Weld 295 322 484 41 65

N = .044% All -Weld 295 387 493 35 80

All -Weld Avg 355 (51.5) 488 (70.8) 38 72

All-Weld 4 724 1217 41 55

All -Weld 4 816 1192 29 33

All -Weld Avg 769 (112) 1205 (175) 35 44

Cross-Weld 295 - 581 (84.3) 66 58

Cross-Weld 4 823 1134 23 27

Cross-Weld 4 792 1109 21 31

Cross-Weld Avg 807 (117) 1121 (163) 22 29

E316L-15 All-Weld 295 336 506 44 69

N = .071% All-Weld 295 385 532 17 11

All-Weld Avg 361 (52.4) 519 (75.3) 30 60

All-Weld 4 829 1236 54 28

All -Weld 4 929 1265 43 49

All -Weld Avg 879 (128) 1251 (182) 48 38

Cross-Weld 295 399 (57.9) 536 (77.8) 17 40

Cross-Weld 4 907 nil 43 26

Cross-Weld 4 910 1263 31 37

Cross-Weld Avg 909 (132) 1187 (172) 37 32

30 a
Table 7. Fracture toughness of shielded metal arc weld deposits

in one-inch-thich 316LN plate

Specimens supplied by the General Electric Company

Temperature

Material K k J/m

(in-i^2)

' in '

Average

<ic(J)

MPa

( ksi/in)

E308L-15 4 102 (582) 156 (142)

4 lOl' (577)

E316L-15

. 044%N

4 58 (331) 118 (107)

E316L-15 4 28.4 (162) 95 (86)

.071%N 4 46.2 (264)

309
Table

8.

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impact

toughness

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weld

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in

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316LN

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77K
REFERENCE

1. Materials Studies for Magnetic Fusion Energy Applications at Low

Temperatures — I (edited by F. R. Fickett and R. P. Reed),

National Bureau of Standards, NBSIR 78-884 (April 1978).

311
TENSILE

YIELD

STRENGTH,

MPa

Figure 1. Influence of Nitrogen on the Strength and Toughness at 77 K of

Three Experimental Filler Metals (SMA weld deposits).

31 L

J
METALLURGICAL FACTORS AFFECTING THE TOUGHNESS

OF 31 6L SMA WELDMENTS AT CRYOGENIC TEMPERATURES

National ELureau of Standards

313

METALLURGICAL FACTORS AFFECTING THE TOUGHNESS OF 31 6L SMA WELDMENTS

AT CRYOGENIC TEMPERATURES*

1 1 2 D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr.

^National Bureau of Standards

Boulder, CO 80303

2Colorado School of Mines

Golden, CO 80401

3

Arcos Corporation

Philadelphia, PA 19143

*Work supported by the U. S. Department of Energy, Office of Fusion

Energy. Contribution of NBS, not subject to copyright.
ABSTRACT

The effects of delta ferrite content, ferrite morphology, carbon

content, and sensitization on the fracture toughness and tensile properties of AWS E316L and E316 shielded metal arc (SMA) weldments at 295,

76, and 4 K were investigated. Ten SMA test welds were evaluated, eight

made with E316L and two with E316 electrodes. In four of the welds, the

delta ferrite content was controlled over the range of 0 to 10.1% ferrite

through slight variations in the chemical compositions of the electrode

coatings. In three of the welds the ferrite morphology was varied in 8%

ferrite welds deposited using the same electrode by varying heat input,

preheat, and interpass temperature. One E316 weld with low ferrite

content (1.2%) and the baseline welding parameters was used to assess

the effect of carbon content. The two remaining welds were made with a

low cooling rate, to study sensitization in the low ferrite, low (E316L)

and regular (E316) carbon welds. Part of the E316 weld was deliberately

sensitized at 675°C for 24 hours.

All of the weldments had excellent toughness at room temperature.

However, at 76 K, only the E316L weld with low ferrite had acceptable

(to ASME Standards) toughness. Large decreases in toughness at 76 and

4 K were related to increasing ferrite content. Decreases in Charpy

impact energy at 76 K were also related to coarsened ferrite morphology

caused by reduced cooling rates, to increased carbon content, and to the

sensitization heat treatment. The tensile-yield strength increased with

ferrite content, especially at 4 K. The effects of ferrite content on

the ultimate tensile strength and the ductility were generally minor.

316
INTRODUCTION

Stainless steel weldments will be used for the main structural

supports of the superconducting magnet systems needed for fusion energy

n 1 devices. The austenitic stainless steels have the best combination

of strength, stiffness and toughness at 4 K, the operating temperature

for superconducting magnets. The service experience with welded assemblies of these alloys has been satisfactory at cryogenic temperatures. A contributing factor to 'this success has been low design

stresses; generally the stresses have been limited to 140 MPa (20 ksi);

the ASME allowable stress for austenitic stainless steels at and below

room temperature.

For large superconducting magnets, the enormous forces that must be

contained and the limited space available for structure may necessitate

operating at significantly higher stresses, stresses based on the strength

of the alloy at the operating temperature. Many of the austenitic

stainless steels exhibit excellent fracture toughness at temperatures

down to 4 K in the base metal. Prior to this investigation, weld metal

fracture toughness had not been determined at 4 K. Charpy impact results

at 76 and 20 K [2, 3, 4, 5] suggested that the toughness may be marginal at

4 K, particularly if the stress levels in the magnets are higher than

current practice.

This paper addresses the question of the effect of delta (6)

ferrite content and morphology on weld toughness at low temperatures.

Delta ferrite, a body-centered-cubic, ferromagnetic phase formed on weld

solidification for a certain range of weld chemistries, is a prominent

feature of the metallurgy of conventional stainless steel weld metals.

From five to ten percent delta ferrite is often introducted to reduce

317
hot cracking, which occurs in many weld metals when ferrite is absent

and renders them unacceptable. While 6 ferrite has not been found

deleterious to the room- temperature toughness of stainless steel weldments, it has been suspected of decreasing low-temperature weld toughness because of the well-known low temperature brittleness of bodycentered-cubic materials such as mild steel. The ferrite morphology is

of interest because isolated patches of a brittle phase should be far

less harmful than a semi-conntinuous network that offers a path for

crack propagation.

Consideration is also given to carbon content, E316L with .03%C vs.

E316 with .06%C, and to sensitization. Again, these factors have a

minor effect on toughness at room and elevated temperatures. Their

possible significance at low temperatures is due to the increase in

austenite strength, approximately two-fold, as temperature is reduced to

4 K. As matrix strength increases, additional sites, such as carbides,

may nucleate voids and reduce toughness.

MATERIALS

Common filler metals for welding austenitic stainless steels for

low temperature service are Types 308L and 31 6L. The 31 6L filler metal

was chosen over 308L for this study because it has higher Charpy impact

toughness at 77 K (-320F) [4,5], and slightly less susceptibility to

microfissuring [6]. The test weldments were 2.5 cm (1 inch) thick

plates of AISI 316 stainless steel welded with either AWS E316L-16 or

E316-16 titania-coated electrodes using the shielded metal arc (SMA)

process. The SMA process was chosen to permit systematic variation of

the weld deposit chemistry by alloy additions to the coating. A singlevee (45° included angle) joint preparation with a 13 mm gap at the root

318
and a 6 mm-thick backing plate was used to avoid dilution by the base

metal at the weld centerline.

Eight 316L and two E316 welds were evaluated. The welding parameters and weld deposit chemistries are given in Tables 1 and 2, respectively. The delta ferrite content was varied from 0 to 10% by altering

the chemistry of the deposited weld metal within the limits for AWS E316L.

The cooling rate of the weld was controlled by varying heat input, preheat, and interpass temperature.- Four E316L weldments made with baseline

welding parameters were used to investigate the effect of ferrite content. Three additional sets of parameters were used to make constant

chemistry E316L welds with varying cooling rates, all with about 8%

ferrite. One E316 weld, with low ferrite content (1.2%) and the baseline

welding parameters was used to assess the effect of carbon content. The

two remaining welds were made with a low cooling rate to study sensitization in the low ferrite welds in both E316L and E316. In order to

further study sensitization, part of the E316 weld (no. 10) was heat

treated at 675°C for 24 hours. This weldment is subsequently designated

number 1 1

.

PROCEDURES

The weldments were evaluated using tensile, Charpy V-notch (CVN)

and fracture toughness tests. Test environments included air at room

temperature (approximately 295 K), liquid nitrogen at 76 K, and liquid

helium at 4.2 K. The procedures are described in the following subsections.

Tensile : Tensile tests were conducted at 295, 76, and 4 K using 2.5 cm

(1 in.) gage length and 0.64 cm (215 in.) diameter specimens. The

specimens were oriented along the length of the weld such that the

319
entire sample was weld metal. The test procedures and the cryostat for

the low temperature tests have been described by Reed [7].

Notch Toughness : Charpy V-notch impact toughness tests were conducted

at 295 and 76 K using the procedures specified in ASTM E23. The specimen length was oriented perpendicular to the length of the weld with

the notch at the weld centerline and perpendicular to the plate surface.

Fracture Toughness : Fracture toughness tests were conducted at 76 and

4 K using the ASTM E399 compact specimen with a 5.1 cm (2 in.) width and

2.5 cm (1 in.) thickness. The specimens were oriented such that loading

was perpendicular to the weld and crack extension was along the weld

centerline. The single specimen J-integral procedure [8] was used to

determine the value of J at the onset of crack extension.

The plane strain fracture toughness, Kj^(J), was determined from

using the relationship proposed by Begley and Landes [9]:

where E is Young's modulus and v is Poisson's ratio. Tests were conducted in an electrohydraul ic fatigue machine equiped with a cryostat

described by Fowlkes and Tobler [10].

In the single-specimen J-integral technique, the specimen is precracked in tension-tension fatigue to a crack length of 3 cm (1.2 in.).

Subsequently, the specimen is loaded monotonical ly at a load-linedisplacement rate 50 pm/s or less under displacement control to a point

well within the elastic region. The curve of load vs. load-line displacement is recorded on an x-y plotter. A specimen is then partially

unloaded, and the incremental compliance (load vs. displacement) curve

320
is recorded on a second x-y plotter with appropriate amplification and

signal conditioning. A signal proportional to the change in load is

electronically subtracted from the signal; the proportionality factor is

adjusted so that the unloading compliance before crack extension is

displayed as a vertical line on the graph of load change against loadline-displacement change. The specimen is incrementally extended, with

several stops for recording the unloading compliance. Typically, ten to

twenty unloading compliances are recorded. The magnitude of the decreas

in load during the unloadings is restricted to ten percent of the total

load. J-integral values are calculated from the load vs. load-linedisplacement record at each unloading point using the Merkl e-Corten [11]

relation. Crack extension values, Aa, are calculated from all unloading

compliance curves except the first one, which is used to evaluate the

compliance at Aa = 0. The (Aa,J) data are then plotted, and J,c is

evaluated graphically as the intersection of a curve through the data

with the theoretical blunting line given by J = 2aAa where o is the

arithmetic mean of the yield strength and the ultimate tensile strength.

RESULTS AND DISCUSSION

Ten test weldments were prepared to study the influence of ferrite

content, ferrite morphology and sensitization on the toughness of stainless steel weld deposits.

Weldments 1 to 4 were welded with electrodes 1 to 4, respectively,

to systematically vary the ferrite content between 0 and 12%. The welddeposit chemistry of each electrode, summarized in Table 2, was within

the chemistry limits specified for AWS E316L. The actual ferrite contents were 0.1, 4.1, 8.5 and 10.1% for weldments 1 to 4, respectively.
Ferrite content was measured with a Magne Gage* using standard procedures [12].

Weldments 3, 6, 7 and 8 were all welded with electrode 3 and the

cooling rate was reduced by increasing the heat input from 6.4 kJ/cm to

20 kJ/cm. The cooling rate in weldment 8 was further reduced by using

preheat and interpass temperatures of 315 ± 28°C (600 ± 50°F); a

107 ± 14®C (225 ± 25°F) interpass temperature was usad for weldments 3,

6, and 7. The idea was to alter the ferrite morphology by reducing the

cooling rate as suggested by Canonico, et. al . [13]. Metal lographic

examination of the welds, Figure 1, revealed a discontinuous ferrite

network in the lowest heat input weld (weldment 6) and a nearly continuous ferrite network in the other three welds. It should be noted

that the welds with 4.1% or less ferrite also had discontinuous ferrite

networks. The higher heat input welds (weldments 7 and 8) had significantly coarser ferrite networks (and effectively, courser grain

sizes) than the lower heat input welds. Cooling rates were calculated

using the relationship suggested by Adams [14].

Weldments 1, 5, 9 and 10, which were nominally ferrite-free were

used to investigate sensitization. Actual ferrite contents ranged from

0.1 to 2.2%. Sensitization was varied by increasing carbon content and

by decreasing cooling rate. Carbon was varied by using electrodes with

low carbon (Electrode 1, AWS E316L, C = 0.03%) for weldments 1 and 9,

and higher carbon (Electrode 5, AWS E316, C = 0.06%) for weldments 5

and 10. Cooling rates were moderate (heat input of 10.1 kJ/cm and a

*This trade name is included for clarity only. No endorsement of this

device by NBS is implied.

322
107 ± 14°C interpass temperature) for weldments 1 and 5, and slow (heat

input = 20 kJ/in and a 315 ± 28°C interpass temperature) for weldments 9

and 10. A deliberate sensitization heat treatment, 675°C for 24 hrs,

was given to a portion of weldment 10. The degree of sensitization in

each of the weldments was not measured. The toughness of each weldment,

which presumably reflects the extent of sensitization, is discussed

below.

Tensile Properties

The tensile properties are summarized in Table 3. The tensile and

yield strengths of weldments 1 to 4 are shown as a function of temperature in Figure 2. The lowest strengths were consistently obtained in

the ferrite-free weld and strength increases of approximately 10% were

obtained in the 8.5 and 10.1% ferrite welds. In the welds made with

electrode 3 (weldments 3, 6, 7, and 8), yield strength increased with

cooling rate as shown in Figure 3. The tensile strength varied ±5%

with cooling rate, but did not increase systematically with cooling

rate. The higher carbon welds (AWS E316, C = .06%) generally had higher

tensile and yield strengths than the corresponding low carbon welds

(AWS E316, C = .03%). Ductility of all the welds was satisfactory with

elongation values ranging from 26 to 51% at room temperature and 20 to

40% at 4 K.

Charpy V-Notch Impact Toughness

The Charpy V-notch (CVN) impact test results are summarized in

Table 4. Note that the results are given in terms of energy absorbed

and lateral expansion. The ASME Boiler and Pressure Vessel Code requires .38 mm (.015 inch) lateral expansion in stainless steel weld

323
metals at the minimum design temperature. Welds made with each electrode and the deliberately sensitized weld easily passed this requirement at room temperature. However, at 76 K, only the AWS E316L weld

with essentially zero ferrite (weldment No. 1) exceeded the .38 mm

requirement. In the subsequent discussion of Charpy test results, the

energy absorbed will be the indicator of toughness because of the excellent correlation of this parameter with the J-integral fracture

toughness as discussed later.

As shown in Figure 4, the CVN toughness values drop sharply as the

ferrite content is increased. The results agree with those previously

reported by Szumachowski and Reid [4]. Comparison of the 76 K toughness

weldments 1 and 5 and weldments 9 and 10 indicates that increasing the

carbon content from .03% (E316L) to .06% (E316) also causes a substantial

reduction in toughness at 76 K. A lesser drop in toughness occurs with

reduced cooling rates. For weldments 3, 6, 7 and 8 lateral expansion

was .16 ± 2 mm and varied linearly with the cooling rate through the

sensitization temperature range, 550 to 800°C, as shown in Figure 5.

Fracture Toughness

Fracture toughness tests were conducted on Weldments 1 to 5 at room

temperature, 76 and 4 K, and the results are summarized in Table 5.

At room temperature, tests indicated that the toughness of each weld

exceeded 275 MPa>7m; specific toughness values were not obtained because

the 2.5 cm-thick specimens did not meet the proposed size requirements [8]

for testing. As shown in Figure 6, the fracture toughness at 76

and 4 K decreased significantly as ferrite increased from 0 to 10%.

These results confirm the trend observed in the 76 K CVN tests and

indicate that toughness continues to decrease as temperature is reduced

324
to 4 K. For the 76 K tests, the Jj^ fracture toughness data correlated

quite well with the CVN data as shown in Figure 7.

Fractography

One fracture toughness test was stopped after extensive sub-critical

crack extension had occurred, but prior to fracturing the specimen. The

specimen was sectioned in the plane perpendicular to the crack surface

and in the direction of crack extension. The crack extension path was

examined with a scanning electron microscope (SEM). The resulting SEM

photomicrograph. Figure 8, indicates that cracking occurs preferentially

along the ferrite phase rather than through the austenite matrix. It

was not clear if the cracking occurred at the ferrite-austenite interface or through the ferrite phase. The fracture surfaces of 76 K

Charpy specimens from Weldments 1 to 5 were examined with the SEM. The

low ferrite weld (Weldment No. 1) had extensive ductile dimpling.

Figure 9. The high ferrite welds had a relatively flat, brittleappearing fracture surface with some regions of ductile dimpling.

Figure 10. Isolated evidence of quasi-cleavage was found, but it cannot

be conclusively stated that cleavage of the ferrite phase causes the

reduced toughness. Visual examination (unmagnified) of the fracture

surfaces of the 76 K Charpy specimens for weldments of each ferrite

level ’ reveal ed increased faceting with increased ferrite content. This

observation coupled with the minimal evidence of cleavage fracture

suggests that cracking progresses along the austenite/ferrite interface

rather than through the ferrite.

The fracture surface of the 76 K Charpy specimen from the deliberately sensitized weld (Weldment 11) was also examined using the SEM.

The fractographs. Figures 11 and 12, reveal that the fracture surface

326
was covered with a precipitate and extensive secondary cracking occurred

perpendicular to the main fracture surface. Apparently, the fracture

path follows the path defined by the localized carbide precipitation

caused by the sensitization treatment.

Comparison with Base Metal Properties

The fracture toughness of the AWS 31 6L weldments at 4 K can be

compared with annealed austenitic stainless steel plate [15]. As shown

in Figure 13, weldments have significantly less fracture toughness than

base metals at comparable yield strengths. This means that the weld

zone is far more susceptible to fracture than the base metal in cryogenic structures. Clearly, further research to improve weld toughness

at cryogenic temperatures is required.

SUMMARY AND CONCLUSIONS

Ten AISI 316 test plates were SMA welded with AWS E316L electrodes

to evaluate the effect of ferrite content, ferrite morphology, carbon

content, and cooling rate on the tensile properties, notch toughness,

and fracture toughness at room and cryogenic temperatures. Each of the

10 test weldments easily passed the ASME Charpy toughness requirement of

0.38 mm (.015 inch) lateral expansion when the test was conducted at

room temperature. However, when tested at 76 K, only the AWS E316L weld

deposit with 0.1% ferrite that was welded using moderate heat input

(10.5 kJ/cm) and interpass temperature (107 ± 14°C) exceeded the ASME

requirement. Increasing the ferrite content to 4.1% or greater, increasing the interpass temperature to 315 ± 28°C, or increasing the

carbon content to 0.06% (E316 instead of E316L) caused the toughness to

fall below .38 mm. The following conclusions were drawn from the test

results:

32C
I 1. Increasing delta ferrite over the range 0 to 10% causes a sharp

decrease in fracture toughness at 76 K and 4 K.

2. The ferrite morphology changes from isolated patches of ferrite

at near zero ferrite to a nearly continuous ferrite phase at 7%

ferrite.

3. Decreased cooling rate increases the ferrite dendritic spacing

and the effective grain size of the austenite matrix, and it decreases the Charpy impact toughness at 76 K.

4. Increasing the carbon content from 0.03 to 0.06% decreased the

fracture toughness at 76 and 4K.

5. Deliberate sensitization of an AWS E316 weld deposit caused a

significant toughness loss at 76 K.

6. Fractography of failed CVN specimens suggests that fracture at

76 K progressed along the ferrite/austenite boundary.

7. The fracture toughness, Kj^(J), of E316L welds is approximately

half of the for AISI 300-series stainless steel plates.

327
ACKNOWLEDGMENTS

This work was sponsored by the U.S. Department of Energy; Dr. E. N. C.

Dalder was the project monitor. The authors wish to thank J. W. Elmer

and Professor D. L. Olson of the Colorado School of mines for metallographic examination of the weldments.

328
REFERENCES

1. H. I. McHenry and R. P. Reed, "Structural Alloys for Superconducting Magnets in Fusion Energy Systems," Nuclear Engineering and

Design, forthcoming.

2. R. Castro and J. J. de Castenet, Welding Metallurgy of Stainless

and Heat-Resisting Steels . Cambridge University Press, 1974,

p. 122.

3. F. W. Bennett and C. P. Dillon, "Impact Strength of Austenitic

Stainless Steel Welds at -320F Effects of Composition, Ferrite

Content, and Heat-Treatment," J. Basis Eng. 8^, 3, March 1966.

4. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic Stainless Steel Weld Metals: Part I - Role of Ferrite,"

Welding Journal, Welding Research Supplement 325-s to 333-s

(1978).

5. W. T. Delong, "Ferrite in Austenitic Stainless Steel Weld Metals,"

Welding Journal, Welding Research Supplement ^ 273-s to 286-s

(1974).

6. C. D. Lundin and D. F. Spond, "The Nature and Morphology of Fissures

in Austenitic Stainless Steel Weld Metals," Welding Journal, Welding

Research Supplement 356-s to 367-s (1976).

7. 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, 1962, pp. 448-454.

329
8 . G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, in

Mechanics of Crack Growth , Proceedings of the 1974 National Symposium on Fracture Mechanics, ASTM STP 590, American Society for

Testing and Materials, Philadelphia, PA (1976), p. 27-42.

9. J. A. Begley and J. D. Landes, 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), 1-20.

10. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a

Ti 6A1-4V Alloy at Liquid Helium Temperatures," Engineering Fracture

Mechanics, Vol. 8, 1976, pp. 487-500.

11. J. G. Merkle and H. T. Corten, "A J-Integral Analysis for the Compact Specimen, Considering Axial Force as well as Bending Effects,"

Journal of Pressure Vessel Technology, Trans. ASME, 1974, pp. 286-292.

12. W. T. DeLong, "Calibration Procedure for Instruments to Measure the

Delta Ferrite Content of Austenitic Stainless Steel Weld,Metal,"

Welding Journal, Vol. 52, 1973, pp. 69-s to 72-s.

13. D. A. Canonico, "Materials and Their Application in the Nuclear

Industry," Metallography, Vol. 9, 1976, pp. 459-481.

14. C. M. Adams, Jr., "Cooling Rates and Peak Temperatures in Fusion

Welding," Welding Journal Vol. 37, 1958, pp. 210-s to 215-s.

15. D. T. Read and R. P. Reed, "Fracture and Strength Properties of Selected Austenitic Stainless Steels at Cryogenic Temperatures , " to

be published.

330
LIST OF TABLES

Table 1

Table 2

Table 3

Table 4

Table 5

AWS E-316L and E-316 Test Weldments.

Chemical Composition of Weld Deposits of the Test Electrodes.

Tensile Properties of the AWS E316L Weldments.

Charpy V-Notch Impact Test Results.

Fracture Toughness of the AWS 316L Weldments at 76 and 4 K.

331
Table

1.

AWS

E-316L

and

E-316

Test

Weldments. o

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(0 CU •>- LO 1 1 S- -M — — CO

fO c ^

D_ t—I o cu CO o o o o o LO cn Ln Ln cn

r— > CM CM CM CM CM CM 1

1— 1r—

CU fO E 3 S- U 1 LO 1 LO 1 Ln 1 LO 1 cn 1 o CD 1 o CD 1 o

[— r— 1— ‘ r— ' CM ' ' ' '

CO

4-> LO LO LO

r— CM CM CM CM CM CM CM CM CM CM o>

-MC • CU CO

O

O

O

O o O

O

OOO S- CL CO CO ro CO ro O LO LO LO LO

s- E n— 1— 1

1

11

1

11

3 C O

O) -oo

s- • +-> o u ^O)

CM ro LO ro ro ro Ln

LU

CU ao _J _1 _J _j _J _i _j _1

s- cu LO LO LO LO LO LO LO LO LO LO

+-> CL (— r— 1— (—

t— 1— I— u 5o CO ro ro ro ro ro ro ro ro CO

CU 1—

LlI

LU LU LU LU LU LU LU LU LU LU

CM ro CO CTi

332

*12

Readings

+Fissures

observed

on

surface

°Weldment

11

was

part

of

Weldment

10

that

was

given a

sensitization

heat

treatment

of

675°C

for

24

hours
Table

2.

Chemical

Composition

of

Weld

Deposits

of

the

Test

Electrodes.

C\J to CM 'd" 2: to •d- o o o o O

JD 'd- to to LO O r (

11

1

LD cn CM 00 LO

O CM CM CO t CM

s: • • • • • CM CM CM CM CM

+-> c I

to LO

Ol •r- <Ti cn o o 00 o z: . • * • • S- CO CM CM 1

CM

cu 1

r— 11—

D_ - -M o to cn

cn S- >=3- to CM to to

•r“ o . • . • . Ol 00 00 cn cn OO 2 '

'

'

'

'

co CO cn CO • r— Q_ ro CO CO CM CO

4-> o o o o o • r— • • . • * toO

Q.E 1

CM cn 00 cn o to (

1— o o o o o O o o o

00 to cn to

OO tn to LO LO LO

CM CM o to LO c

s: • • • • • o ca- 00 <n o CO CO CO CM LO o o o o o

"O . (U o S- tn to o o to 2: =3 • • * « * to o CO cn CM o O) fO 1

+-> <u

•1

2:

sS-O) E

Li_ 03

S- to LO o cn • • Qi 03 CM o CM CO 2 • 1

r— 1

Q

CU T3O

S- • +J o 1— CM CO 'd" LO u z:

(V n—

UJ

LD

4-> (OOU

OJQ.

l/)

t/5

SQJCL

-M

l/) E

o»u

-!->

LOO

CL

O) •O

"O

OJ

333
Table 3. Tensile Properties of the AWS E316L Weldments.

Weldment

No.

Ferri te

%

Temp.

K

Yield Strength,

0.2% Offset

MPa(ksi

)

Ultimate Tensile

Strength

MPa(ksi

)

Elongation

in 4-d

%

Reduction

of Area

%

1 .13 295 392 (56.9) 535 (77.6) 38 58

II 76 659 (95.6) 983 (143) 19 28

II 4 694 (101) 1222 (177) 20 28

2 4.1 295 396 (57.5) 552 (80.1) 35 57

II 76 657 (95.3) 1064 (154) 30 32

M 4 798 (116) 1331 (193) 38 42

3 8.5 295 452 (65.6) 578 (83.9) 45 63

II 76 753 (109) 1183 (172) 20 42

II 4 886 (129) 1294 (188) 28 32

4 10.1 295 429 (62.3) 590 (85.6) 38 66

II 76 725 (105) 1203 (175) 40 32

II 4 881 (128) 1373 (199) 38 47

5* 1.2 295 405 (58.8) 555 (80.6) 32 56

II 76 680 (98.7) 914 (133) 12 24

II 4 924 (134) 1364 (198) 32 40

6 6.5 295 509 (73.9) 609 (88.4) 26 52

II 76 793 (115) 1261 (183) 55 34

II 4 936 (136) 1325 (192) 36 32

7 9.7 295 441 (64.0) 564 (81.9) 46 64

II 76 736 (107) 1179 (171) 58 48

II 4 817 (119) 1369 (199) 39 54

8 7.3 295 412 (59.8) 585 (84.9) 51 51

II 76 678 (98.4) 1175 (171) 57 48

II 4 792 (115) 1263 (183) 40 46

9 .67 295 343 (49.5) 514 (74.6) 48 59

II 76 639 (92.7) 996 (145) 30 30

II 4 745 (108) 1157 (168) 26 26

10* 2.2 295 417 (60.5) 597 (86.6) 44 55

II 76 711 (103) 1250 (181 ) 46 29

II 4 847 (123) 1158 (168) 26 26

*AWS E316 electrode.

334
Table

4.

Charpy

V-Notch

Impact

Test

Results.

c CTi CO CO CO o CO CO CO CO o c: LD CO r-s. CD CO LD LO CM CO CO

r— T— 1— r— o o O o o o o o 00 O o o o O o o o o o o c

03QLlJ UD

I

I

,1 n

O C\J V£) oo

OO (— C\J I r—

C\J oo

cri o

CD +-> S- MUD CUc

IX)

C\J

UD 00

C\J CD

CD

C\J

C\J CO

C\J CO

ro

C\J CM oCM

(1) Se LO CD un o CO o e CO CO LD LO LT> LT) LO CO CO LO

•r— o o o o o o o o o o o OOc

fT3 Q_

X

LlJ

13 fO +J S- n3 CD

S- +j

OJ <T3 E

Q-_| E

«;d-cDCMroc^CM'^cor^Ln«^ ocn^oor^CMD-ouDun

O)

o

o

QC JD

>>CD 4-> S- MOJ

"-D

CO CO CM CO CO CO CM LO CO o

(/) -M CO CO LO CO CO CO CO CO CO LO CO

o CO CO CO CM CO CM CM

CU

+-)

S-o

CD o CO CO CD CO CO CD CO o

LO CM LO CO CM 1

C|-

i-o

4-

1 to o CO o r— CO CD o CM 1 (U

« 13

Ol -Oo

S-o

O)

LlJ

CM ro ^ LD CO CO CO un LT)

fO>

(UCD

fO

s-OJ>

fO

O)

S-03

03

c

OlE • o o

I— 2:

(U

CO uo CO

03 Q

CD CD

4-> o

335
Table

5.

Fracture

Toughness

of

the

AWS

316L

Weldments

at

76

and

4

K.

COcn ocri u o

00

Oni

00

CM

CD oo

CO

CM OCO ocr>

u ^ I—I fO D.

(OOCM<Ti o

CO

CO COo

00

<T> 00

CT> OO

LD ID

CM

JOc

o o CM<— 00

O CM

CD CM

CTi LD

Q.

E

oj CD CD

LD on

cr> CM 1

Cr> CD CTi

0)cn

fO

s-<u>

<c

CD

on o O CD

'Cf

o

r-^ oCM LO

00

LO

CO

cucn

(T3 s-O)>c

CD O CM lO

CD CD

00 00 cr.

CX) CD CTi CO LO

00

CO

cucn

n3scu>

<c

CD

LO CD O

'cf *=dCO CD CM

LO ^

LO

00

CUcn

(O

s-cu>

<c

CTi cn

CD CD

cn CD cn

CM I

CD

00

cucn

fC

S-

(U>c

CD

^ CO

'=3'

CM

COCM

CM cn

LO CO

CO I—

CM ^ CO CM

CUcn

fO

Scu>c

CD CD

CD CD cn

CM CM I—

CO

•cn

cucn

(O

cu>

«a:

CD

CO

CM

CO

CM ^ LO

LO CD

LO 00

cn I—

cucn

<o

s-CU>

<c

CO CO

CD O CO O CD

CM

00LO

(U +J

•r-' ss-cu

Ll_ 5^

LO

00

CM

+-> c

cuE

-o

"oj

CM CO

JtLO

336

76

Average

107
LIST OF FIGURES

Figure 1

.

Photomicrographs showing the morphology of 6-ferrite in an

austenitic matrix in 31 6L weld metal deposited by the SMA

process at four heat input levels.

Figure 2. Temperature dependence of the yield and tensile strengths of

E316L weldments with 0 to 10% ferrite.

Figure 3. Cooling rate dependence of the yield strengths of E316L

weldments with 8 ± -2% ferrite at room temperature, 76 and

4 K.

Figure 4. Charpy V-notch impact energy at 76 K as a function of ferrite

content for E316L weldments.

Figure 5. Charpy V-notch impact energy at 76 K as a function of cooling

rate for E316L weldments with 8 ± 2% ferrite.

Figure 6. Fracture toughness, at 76 and 4 K as a function of

ferrite content for E316L weldments.

Figure 7. Correlation of the critical J-integral, Jj^, with Charpy

V-notch impact energy for E316L weldments at 76 K.

Figure 8. SEM photomicrograph showing the crack propagation path in a

31 6L weldment with 10% ferrite (2000X, reduced 55% for re

production)

.

Figure 9. SEM photomicrograph of the fracture surface of a 76 K

Charpy V-notch impact specimen from an E316L weldment with

0.1% ferrite (550X).

Figure 10. SEM photomicrograph of the fracture surface of a 76 K

Charpy V-notch impact specimen from an E316L weldment with

10% ferrite (550X).

Figure 1 1

.

SEM photomicrograph of the fracture surface of a 76 K

Charpy V impact specimen from an E316 weldment that was

deliberately sensitized at 675°C for 24 hours (550X).

337
Figure 12. SEM photomicrograph of the fracture surface of a 76 K

Charpy V impact specimen from an E316 weldment that was

deliberately sensitized at 675°C for 24 hours ((a) 2200X and

(b) 5500X).

Figure 13. Fracture toughness at 4 K of E316L weldments as a function

of yield strength. Also shown is a best fit line for the

stainless steel plate of fracture toughness in yield strength.

338
Figure 1. Photomicrographs showing the morphology of del ta-ferri te in an austenitic mat>^i\ in 31 6L weld metal deposited by the SMA

process at four heat input levels:: 5.3, 10.5, 13.1, and 19.6 kJ/cm

in photomicrographs (a), (b), (c), and (d), respectively. 330 X.

339
Strength

(MPa)

leoor T T

0 J L.

100 200 Temperature (K)

300

Figure 2,. Temperature dependence of the yield and tensile strengths of

E316L weldments with 0 to 10^ ferrite.

340

Strength

(KSi)
YIELD

STRENGTH,

MPa CO

(S>

cc

I—

ooQ

>-

Figure 3. Cooling rate dependence of the yield strengths of E316L

weldments with 8 ± 2% ferrite at room temperature, 76 and

4 K.

341
Figure 4. Charpy V-notch impact energy at 76 K as a function of ferrite

content for E316L weldments.

342
Cooling Rate At Sensitization (°c/sec)

Figure 5. Charpy V-notch impact energy at

rate for E316L weldments with 8

76 K as a function of cooling

± 2% ferrite.

343
Figure 6. Fracture toughness, Kj^(J) at 76 and 4 K as a function of

ferrite content for E316L weldments.

344
Figure 7. Correlation of the critical J-integral,

V-notch impact energy for E316L weldments at

with Charpy

76 K.

345
Figure

8.

SEM

photomicrograph

showing

the

crack

propagation

path

in

31

6L

weldment

with

10%

ferrite

(2000X,

reduced

55%

for

reproduction)

.
Figure 9. SEM photomicrograph of the fracture surface of a 76 K Charpy V-notch impact specimen from an E316L weldment with

0.1% ferrite (550X).

347
Fiqure 10. SEM photomicrograph of the fracture surface of a 76 K Charpy V-notch impact specimen from an E316L weldment with

10" ferrite (550X).
Figure 11. SEP. photomicrograph of the fracture surface of a 76 K Charpy V impact specimen from an E316 weldment that was deliberately sensitized at 675°C for 24 hours (550X).

349
(a)

(b)

Figure 12. SEM photomicrograph of the fracture surface of a 76 K Charpy V impact specimen from an E316 weldment that was deliberately sensitized at 675°C for 24 hours ((a) 2200X and

(b) 5500X).
YIELD STRENGTH (ksi

)

0 100 200

I

1 1 1

1 r

YIELD STRENGTH (MPa)

Figure 13. Fracture

of yield

stainless

toughness at 4 K of E316L weldments

strength. Also shown is a best fit

steel plate of fracture toughness i

as a function

line for the

n yield strength.

3S1

FRACTURE

TOUGHNESS

(ksi/Tn)

MICROSTRUCTURE OF SPECIFIC STAINLESS STEEL WELDMENTS

BEING EVALUATED FOR CRYOGENIC SERVICE

Colorado School of Mines

FINAL REPORT

"MICROSTRUCTURE OF SPECIFIC STAINLESS STEEL

WELDMENTS BEING EVALUATED FOR CRYOGENIC SERVICE"

J. W. Elmer, P. A. Steinmeyer and D. L. Olson

Department of Metallurgical Engineering

Colorado School of Mines

Golden, Colorado 80401

Submitted to:

National Bureau of Standards

Cryogenic Laboratory

Boulder, Colorado

November 28, 1978

355
INTRODUCTION

Stainless steels derive their corrosion resistance from the presence

of specific alloying elements; the major addition is chromium which,

when present in amounts in excess of 11%, foms a protective chromiumoxide layer on the surface of these alloys in corrosive environments.

Other alloying elements added to stainless steels are Ni, Mo, Mn, Ti, Nh,

Cu, and N. Each influences the microstructure in a specific manner.

The microstructure of stainless steel greatly influences its mechanical

(1-3) and corrosion behavior. Welding will change its microstructure to

such an extent that the weld will have properties different from either

the base or the filler metal. During the welding process on most stainless

steels, the molten metal solidifies into a microstructure with two phases;

austenite and/or ferrite''*^ . The relative amounts of each structure depend

on several variables such as welding heat input welding technique^^^

f r\ i O 7 and weld metal composition'' . Martensitic structure, with a c/a

ratio approaching one may also form with mechanical working or with an

extremely low quenching temperature.

Much work has been performed to relate the microstructure of the weld

metal to its composition. The first of such work was published in 1949

( 8 ) by Schaeffler' with his constitution diagram for stainless steel welding

(Figure 1). The Schaeffler diagram provided a means to determine the weld

deposit ferrite content by combining the effects of the austenitic stabilizing elements Ni, C, and Mn with the ferrite stabilizing elements Cr,

Mo, Si, and Nb.

( 12 ) Seferian has also proposed a mathematical expression for ferrite

content. His equation is derived from the Schaeffler diagram, and allows

the calculation of ferrite content of an austenitic stainless steel weld
C«2a§;l£!liiiil£<©ira

ff®r

§toSnnI!©§s

Steel

WeM

Metol

ByA

nton

L,

SchacJJJer

3G7

Figure

1:

The

Schaeffler

diagram

for

stainless

steels

which

provides a

means

to

measure

weld

deposit

Ferrite

content

based

on

weldment

composition.
deposit as follows:

% ferrite = 3{Cr - .93 Ni - 6.7}

eq eq

(9) In 1956, DeLong expanded on the Shaeffler diagram to include the strong

austenitizing effects of nitrogen (Figure 2).

Recently Hull^^^^ has modified the DeLong diagram to include the

inversion behavior of manganese from an austenitic to a ferritic stabilizer

at high manganese levels (4% of larger) . The Hull diagram is compared

with these other diagrams in Figure 3. Notice the various expressions

for the nickel and chromium equivalents. The tendency for other alloying

elements to act as austenite or ferrite stabilizers has also been studied.

For example, the presence of copper has been found to Increase austenite

stability in stainless steels, although its effectiveness is only about

0.6 times that of nickel. Thus, in the expression for nickel equivalent,

copper has a coefficient of 0.6. Similarly, the effect of tungsten (a

ferrite stabilizer) is accounted for in the chromium equivalent expression

by assigning it a coefficient of 0.5, while titanium has been reported to

have a coefficient varying from 2 to 5.

In recent years the role of delta ferrite has become an important

(1 30-33) factor in explaining the behavior of weldments ’

. Some of the

general effects that delta ferrite has on weldments are:

1. The ability to control hot cracking at 3 to 8% ferrite.

2. An increase in tensile strength with higher ferrite contents.

3.

Reduction of stress corrosion cracking at high ferrite levels (57)

At high temperatures (1000 F) the delta ferrite transforms to the

more brittle sigma phase reducing ductility, strength and corrosion

f (41,58,59) resistance of the steel

4 .
The

DeLong

Constitution

Diagram

359

Figure

2.

The

DeLong

diagram.

This

is

based

on

the

Schaeffler

diagram,

but

is

revised

to

include

the

strong

austenitizing

effect

of

nitrogen.

The

boxed-in

area

indicates

the

range

of

compositions

allowed

for

31

6L

stainless

steel.

The

heavy

line

sloping

upward

from

left

to

right

is

the

liquidus

line.
A.

B.

C.

4

c

•I S

c

®

o

>

*3

cr

in

4* o

liJ §5

O

Z ro

Cr. Equivalent = ( 7© Cr 4 ®/o Mo + 1.5 x%Si 4.05x7©Cb)

Cr. Equivalent = (1.21 x 7©Mo + 7© Cr 4.487©Si +.l47©Cb)

Figure 3. Comparison of the three constitution diagrams.

A) Schaeffler Diagram, B) DeLong Diagram, C) Hull Diagram

360
5. Increases the susceptability of stainless steel weldments to

porosity.

Lundin and Spond^^^ have reported on the nature and morphology of the

fissures found in stainless steel with low ferrite levels. They found

that fissures occur primarily along grain boundaries in the heat affected

zone of the previously deposited pass. The fissuring tendency is enhanced

by multiple thermal cycling in the heat affected zone and is invariably

more pronounced in ferrite-free areas. They also found that the fissures

form by a liquation mechanism. Average fissure length was found to be

0.004 in. (0.1 mm) in stainless steel weldments. These fissures were

found to be resistant to propagation at room temperature even at strains

of up to 20%.

Sulfur and phosphorus tend to promote liquation in stainless steel,

/

3

^

)

and may result in fissuring . Sulfur has been reported by Bernstein

(26)

et al'’ to have more influence than phosphorus on mechanical properties;

sulfur level should be kept below 0.015% for critical applications.

(37) Honeycombe and Gooch have found that manganese additions reduce cracking

but also must be controlled, as manganese may have deleterious effects

on corrosion resistance. They showed that additions between 2% and 6%

manganese produced crack-free welds.

/ or N Brooks and Lambert'' also report that the amount of ferrite necessary

to prevent cracking is directly related to the impurity level (P + S).

They report that stainless steel welds with ferrite contents greater than

ferrite number 13 to 14 were able to accommodate large amounts of phosphorus

f 38^

and sulfur without cracking. Haddrill and Baker' have reported that

increasing the carbon content apparently reduces the cracking tendency

in austenitic stainless steel. Furthermore, the use of a large-diameter

electrode decreases the fissuring tendency.

3C1
The presence of controlled amounts of delta ferrite in austenitic

stainless steel weldments, although beneficial in reducing microf issuring

,

may have an adverse effect on mechanical properties. The influences of delta

ferrite on high-temperature mechanical properties have been studied by a

number of investigators^^^’^^ ‘^5). et al^^^^ reported that anisotropic

deformation behavior occurs that can be related to the orientation of the

solidification microstructure of the weld. Transverse weld metal specimens

were found to be consistently stronger than longitudinal specimens. During

creep testing the original ferrite transforms to sigma phase. The resulting

hard precipitates of sigma phase were aligned with the solidification substructure making high temperature deformation anisotropic. Type 316 stainless

steel welds with controlled amounts of residual elements (CRE) such as boron,

phosphorous, and titanium have shown good strength and ductility at 649C in

both short-term tensile tests and creep tests lasting up to about 6000 hr.

(43 46) Shahinian et al ’ evaluated the fatigue behavior of stainless steel

weldments at 25°C, 427°C, and 593°C, and found type 316 stainless steel weld

metal to have better resistance to fatigue crack growth than type 308 stainless steel weld metal. They also found that 304 stainless steel weld metal

(33) had better resistance than 316 stainless steel base metal. Thomas reported

that optimum creep resistance at 600C occurs at approximately 5% ferrite

(33) content. Thomas also found that at high ferrite levels ('v<10% ferrite),

the ferrite forms a continuous network, significantly increasing creep

resistance

.

Charpy V-notch and J data (as determined from single-specimen J

integral tests) have shown a marked decrease in toughness with increasing

ferrite content. When the amount of ferrite was increased from 0.13% to

10.1%, the 76K CVN energy dropped from 35. 2J to 12, 6J, while

3G2
2 2 the 76K fracture toughness decreased from 174KJ/m to 43.0KJ/m . Fracture

2 toughness tests at 4K showed the same trend, dropping from 135kj/m to

2 4lKJ/m . Tensile tests showed only slight increases in yield strength at

higher ferrite levels.

The loss of toughness at high ferrite contents in austenitic stainless

steel weld-metals is believed to be due to the brittle behavior of the ferrite

at lower temperatures. Transmission and scanning electron fractography indicates

that high ferrite samples fracture by a cleavage mechanism, while low ferrite

fracture is characterized by microvoid coalescence. The SEM micrograph of

Figure 4 depicts a crack in a 4K single specimen fracture-toughness test sample;

the specimen has been partially cracked, unloaded, and sectioned. The dark

areas in the figure are delta ferrite dendrites. It is seen that the ferrite

influences the crack propagation behavior: cracking occurs more readily along

ferrite regions than through the more ductile austenite matrix.

(47) Ekerot et al have found that the nature and number of pores formed

in austenitic stainless steel weldments can also be related to the stainless

steel microstructure. They reported that primary precipitation of pores

occurs on ferrite dendrites.

There are four major factors which control the amount and morphology

(61) of ferrite in the weldment. These are heat input , amount of applied

strain, composition and history . Williams reported that higher

peak temperatures resulted in higher ferrite contents. Also, the applied

strain during the thermal cycle of welding has been reported to influence the

amount of ferrite. These results suggest that the size of the weld deposit,

which has an effect on the amount of shrinkage, is likely to influence the

amount and morphology of the ferrite.

A number of investigators have evaluated the influenc^e

of the electrode coatings on the amount of ferrite formed

363
364

Figure

4.

Crack

tip

region

of

compact

tensile

specimen

sectioned

from

316L

weldment

(Ffl

=

11),

shovying

crack

propagation

behavior

affected

by

presence

of

delta

ferrite.

Fracture

at

4K,

SEil

micrograph,

2000x.

Reduced

55%

for

reproduction

.
studied weldments (25) and solidification cracking resistance. Binkley et al

made with 308 stainless steel electrodes with lime, lime-titania and titania

based flux and found lime-covered electrode deposits to be weaker in creep

rupture testing at 1200F than either lime-titania or titania covered electrode deposits. Daemen^^^^ and Zvanut and Farmer^^^^ have also been .r

developing stainless steel electrode materials. „ . „ (15,34,49,50,51,62,63,64) ^ ^ . Many investigators » » » > / performed research

to determine ferrite levels needed to control hot cracking. It is wellknown that hot cracking may be eliminated by residual ferrite contents as

small as 2 to 3%.

Recent work has shown that hot cracking is not only a function of the

residual ferrite content, but of the primary phase to solidify; if the

primary phase is austenite, then fissuring occurs, but if the primary phase

is ferrite, no cracks form. The explanation lies in the fact that sulfur

and phosphorus are retained in the primary ferrite, thereby controlling

liquation if ferrite forms first. But if austenite is the primary phase,

then sulfur and phosphorus are left in the interdendritic liquid where

liquation can occur. Thomas explains the role of the primary dendritic

phase by studying the liquldius diagram for the iron-chroraium-nickel ternary

system (Figure 5) and looking at a constant iron isopleth (Figure 6). From

these diagrams, weldments can be classified as having primary ferritic or

austenitic dendrites. This means that weld metals on the austenite side

of the liquidus line are susceptable to hot cracking whereas weld metal

on the ferrite side does not promote liquation along the dendrite boundaries.

Because of the importance of delta ferrite in weldments, a number

of devices have been developed to measure residual ferrite content. These

Instruments - the Severn gauge, Magne-gauge, and Forrester ferrite content

meters are all based on the same principle—the measurement of the magnitude

365
Cr

Figure 5: Iron-Chromlum-Nickel ternary phase j equillbra' diagram.

|

1

!

I

36C
I

3C7

Figure

6.

The

four

different

solidification

structures

v;hic'h

occur

at

the

indicated

compositions.

A

constant

percent

iron

is

assumed

and

the

structures

are

drav;n

at

a

temperature

directly

bellow

their

respective

solidus,

before

appreciable

t

r.ansf

ormat

ion

has

taken

plac'e.
of magnetic attraction in the steel. This attraction is due to the ferromagnetic nature of the ferrite. These ferrite meters measure the volume

percent ferrite in the steel. However, due to different calibration methods

and interest differences associated with each type of meter, accurate

ferrite readings have been difficult to obtain. To solve this problem,

a standardized calibration technique has been agreed upon by the welding

(52) community . The method agreed upon was to replace the percent ferrite

with the ferrite number (FN), where the ferrite number is a magnetic response

obtained from test blocks with various non-magnetic coating thicknesses

produced by the National Bureau of Standards.

Goodwin et al^^^ have shown that variation in ferrite distribution

is related to weld heat input; this variation can cause difficulty with

(25) ferrite content measurement. Binkley et al also have shown that even

variation within AWS compositional specification limits produced significant

variation in mechanical properties. Although the ferrite readings taken

using the ferrite number give a more accurate means of comparing data, the

problem still exists with reproducing a weld with the same ferrite level

because the retained ferrite is not only a function of composition but also

a function of the selection of the welding process and parameters

(49) Lindin et al determined the minimum ferrite contents necessary

for the elimination of fissures in several weld metals. They found that

316L stainless steel weld metal required a ferrite number (FN) of 1.5,

while 316 stainless steel required an FN of 2.5. While it is known that

at least 2% delta ferrite is necessary for elimination of microf issuring,

there still is little information as to the mechanical behavior when the

delta ferrite content is above 10%. Above 10% ferrite, the delta ferrite

may be continuous in the solidification direction. These weldments with

363
the continuous delta ferrite phase will most likely pass through the ductilebrittle transition on cooling and have very limited mechanical properties.

The purpose of this investigation was to assist in the characterization

of the microstructure of a set of stainless steel weldments which have

been extensively mechanically tested for cryogenic service. Each weldment

has experienced either a different filler metal or welding procedure.

Experimental Procedure

Materials and Test Matrix

The weldments were prepared using a one inch, 316L stainless steel

plate; all welds were multipass, covered electrode butt welds. Five electrodes were used. Each was designed to give a different ferrite level in the

weldment. The compositions of the electrodes and base metal are listed

in Table I. The experimental test matrix is separated into two sets of

five weldments, the first five weldments being welded with constant welding

parameters but using each of the five separate electrodes; the second set

of five weldments was welded with select electrodes varying the welding

parameters to produce weldments with differing energy inputs. The welding

parameters for all ten weldments are listed on Table II.

Weldment Preparation

In order to determine information about the welds before they were

sectioned, a nitric-sulfuric acid macro etch was prepared and the as received

weldments were placed in the boiling etch. The weldments were observed

periodically over a 72 hr. exposure to the etch. This preliminary procedure

gave little microstructural information, but was of assistance in the

determination of proper sectioning. The weldments were then sectioned,

ground, and polished. To bring out the microstructure a 10% oxalic acid

3G9
Held

fiGtal

Composition + CN + 1—1

in CN o ro 0 c • • • • • o Pm CTi 1—!

-p 1—1

•nu ro u

0) ++ 1—1• 1 1• m• 1—1• CN o Pm o\o 00 o 1 1 1—1 + +J 3 0

•H rH Cr in o LO g 0 0 CO 00 00 <y\ 0 > U • • • • U 1—I 1—1

CN CN 1—I ^ CN CN CN CN CN CN U w

-p + c

rH 0 0 H Cr CN

H CTi CN in

M fd 0 O 1 1 1 1 00 in o > •H • • • • • •H *H in iD S ;3 —1 1—1 1—1 1—1 1 1 cr w

f'O on ro 1 ro ro CN ro 1 (h o t o O O 1 • o • • • 1

rH CN CTi 00 CTi 1 1—1 1 1 O o O 1 w o O O o O 1 * • • • • 1

in in in 1 Xi fH 1 1 rH 1—1 1—1 1 u • • • • • 1

00 00 C3> VD r'

H in in in in in CN

o\o

CO • • • • •

in CT\ CN 00 in o c 0 CN CN ro 1—1

CN o 0 S • • • • • • •Hp

CN CN CN CN CN ro

•Hw

O iO <y\ 0 l£) CN iO \D in

04

u • • • • • B u 00 CO a^ cr> 00 00 0u I—1 1 1 1—1 1—1 rH 1—1

CN CN o iO in 1—1 c r' r' s • • • « • • 1—1 1—1 r-H I—1 1—1 1—1

CN lO CN ID CN • O o O o o O • • • • •

rH ID in

CTi (Ti o O 00 00 •H • • • • • • s ro CN CN 1 1 CN CN 1—1 1 1 1 1 1 1 1 1 rH 1—1 O 00 CTi 00 u

ro ro ro CN in o • o o o o ! o • • 0o 0 p

-l-> PI ro •U 1—1

CN ro in KO <— o H ^

Q)

r-^ 0 w

CO

<T3 CD 370 .

Using

DeLong

equation

(Reference

9)
Welding

Parameters P c P 3 -H

O rH <T

O CT \ (0 D4 • • • • • • • 1 CD C -r-i vD ID vD ID ID ID vD cr« cn CTi ac H a; CN CN CN CN CN P 'S' 'T 1

C/l Ip (D 0 CO rH

m CTm ro CT vD CN

m 1 CO fO CN CN m m ro CN rH ni CN 1 4fc fO Oh

x3

a

<H 6 ti in

mm

mm

mm o o o 1 0) (1) 0 CN CN CN CN CN CN CN o o o 1 3 E-i CM CN CN CN CN CN CN VD vD VD 1

0)

4-> X3

C

<u CD •;! 0) E

a, > o 1 e

(0 OD1 CD1 00 CD1 001 rH1 VD1 vD1 VD1 vD1 11

fO cn c vx ID ID >X) ID 00 •vl' 'T 1 04

CO c p 5!j VD vD vD 1 •H rH CN CN CN CN CN CN CN CN CN CN 1 X! 0>

QJ 3 pc ^

QJ CO o O o o O

O o o O

O 1 P cu m ro ro on ro o VD VD VD VD 1 e 3 < U

rH rH rH rH rH rH rH rH rH P 1

H- CM cn rH in CN o CO rH in CT

0)C O' • • • • • • • • • O -p &H > o (T rH rH o r- CN •H < rH rH pp

Cl)

m

CiL^ rH rH

m rH CN in p' CO vD CN

IT' • • • * • • « • • O dP > o 00 o rH ID CT p' pvl < rH

0) X)0 P P 0)

4-)

p XI rH CN ro

mm ro ro rH in 1 c <-> ^

0) 0) 3 e P 2

Xi

W

rH

0) 3

oi cc Oi cd oc; 0) P in ID CO CT p^ in cr> P 0) o O o o o O M O H o > o > o >

03

n o o o o o O

M

OM

OH

O

M

OH

J

-H E (N CN CN CN CN CN 1 CN

M CN 1 CN 1 PN 1 VD

Oh

D rH rH rH rH rH P cc

p 1 P 03 P 03 rH 2 E- Eh Eh Eh Eh Eh 03 Eh Eh Eh ro

371

Magna-gage

measurements
electroetch was used operating at a current of 2 amp. for 30 secs. This

etch worked extremely well in delineating the fusion zone between individual

weld beads, as well as the solidification structure. The oxalic acid etch

was used throughout the remainder of the experiment because of its etching

ability and because of the control of the etching parameters. Metallographic procedures for stainless steel weldments are given by White and

LeMay^^^\

Macro and Microstructure

After the welds had been sectioned, polished and etched, a macro

photograph was taken of each weld at a 1.5x magnification. To compare

the microstructure of each weldment a series of ten photomicrographs was

taken across the weld bead at lOOx magnification. The welds were then

examined from a direction looking down on the top of the weld bead by

sectioning along the weldment. This section was then mounted in bakelite,

2

polished, and etched in oxalic acid. A light oxalic acid etch (.5 amp/cm

for 20 sec.) was used in order to bring out the microstructure of the delta

ferrite. The Forrester ferrite content meter was used on each weldment

to measure the percent of delta ferrite. This instrument allows measurement

of the ferrite in each bead of the weldment. It was found that weld beads

which had melted the parent metal were usually lower in ferrite content

than beads in the interior of the weldment. The averaged measured values

agreed well with the Magne-gauge measured values given in Table I. The

microstructure was thoroughly characterized using both optical and scanning

electron microscopy.

Two ASTl'I corrosion tests were performed on the weldments to bring out

better microstructural characterization of these weldments; the Streicher

oxalic acid test and the Huey nitric acid test. The Streicher test was

372
performed on the polished sample, which included the entire fusion zone,

heat affected zone and base metal. The electroetching was performed in an

electropolisher with a constant size etching aperture; the current

density was 1 amp. per sq. cm. of aperture. The Huey test was performed

on two identical sets of samples, corroded in separate vessels so that a

means existed to verify the data taken from each vessel. The two samples

were prepared by sectioning them out of the center of the fusion zone using

an abrasive cutoff wheel. The samples were then boiled in a 65 weight

percent nitric acid solution for five consecutive periods of 48 hours each.

At the end of each time period the samples were removed, cleaned, dried

and accurately weighed to determine the weight loss.

Results and Discussion

Macrostructure

In the polished and etched weldments, the individual beads are readily

distinguishable in the fusion zone as shown in a representative macrograph

of one of the weldments (Figure 7). In this macrograph, epitaxial growth

is apparent, and the grains seem to grow in a preferred orientation which

is perpendicular to the fusion front. It is also apparent that the grains

not only grow in a preferred orientation, but they also grow in a sequential

manner from bead to bead, thus giving the weldment the appearance of having

large columnar grains growing in from the base metal towards the top center

of the fusion zone. It is expected that such an orientation of the solidification structure will strongly affect mechanical properties. This suggests

that special orientation is necessary for the extraction of mechanical

specimens.

373
Figure 7. ilacro Photograph. 1 . 5x mag. 07-IV Rod 3. Showing

individual beads in the fusion zone and epitaxial

growth in a large columnar type structure.

574
Microstructure

A series of microphotographs was taken across each of the etched

weldments to compare the microstructures; a representative micrograph is

shown in Figure 8. Several observations can be made by comparing these

photomicrographs. There is an apparent decrease in dendrite arm spacing

with increasing ferrite content. This decrease in dendrite arm spacing

can be attributed to a large number of ferritic dendrites at high ferrite

compositions which have nucleated and grown. This reduction in dendrite

arm spacing has little or no effect on the sequential grain growth from

bead to bead, as was found by comparing the microstructure of the high and

low percent ferrite weldments. Such sequential growth suggests a weakness

in the weld by allowing crack propagation across individual beads. Another

observation was the large variation in the nature of the microstructure

at different positions across the weldment.

Delta Ferrite

In an attempt to determine the solidification nature of the ferrite

the original weld samples were sectioned in a direction perpendicular to

the top plane of the weld bead and etched to show the ferritic structure.

The micrographs of Figure 9 illustrate the appearance of the ferrite structure at three levels of ferrite content. Notice the increasing continuous

ferrite network with increasing ferrite content. Figure 9C contains A.1%

ferrite (4.5 FN) and can be considered only semicontinuous when compared

with the structures of Figure 10 which range from 6.5% (7.1 FN) to 9.7%

(10.5 FN) ferrite, all containing a nearly continuous ferrite network.

Figure 9A provides an insight into the nature of the solidification

mechanism of the stainless weld metal. In this picture the dark areas are

retained ferrite. Surrounding the ferrite is a light gray phase which is

375
37G

Figure

8.

Sequential

grain

growth

from

bead

to

bead

in

the

weldment.
o OJ

s_

OJ1

CD

S-oS+-> • <U -£= C O

4-> CO O)

=3

O

O Sr3 4-> c u

•r- O)

-4-> 1—

£Z OJ oO "O

•r“ fO oo

4-o o C I— O fO

•iX

-M

O 2C X o o M- OL.") OJ 23 +Jc

CD CJ C 4-> •r- C 3 o o u

CO OJ

CO -f— ^ s- O- s- fO <v

S- M- CDO

s-u

CD

CO

fOOJ

S-o

CD

OJ

S-3

CD

CD

S-s_CD

COO

LD

-ao

ccI CDO

C\J 1

C\J 'J- o

T3 "O O O

CC OC >

(—< I—I I I CD CO O O

377

ferr i

te
the austenite that transformed from the original primary ferrite as the

stainless steel cooled below the solidus temperature. The remaining dark

gray area is the austenite resulting from the transformation in ferriticaustenitic eutectic structure which solidified directly from the molten

steel. This solidification pattern shown in Figure 9A is predicted from

the isopleth of the iron-chromium-nickel ternary phase equilibrium of

Figure 5.

Figure 10 shows micrographs of weldments made using the same electrode,

but welded with varying welding parameters. The microstructure which

results from increasing the energy (heat) input into the weld is illustrated in Figure lOA to Figure lOD as indicated. The ferrite content

varies from 6.5% (7.1 FN) to 9.7% (10.5 FN), but was not found directly

related to heat input suggesting that the variation in ferrite is due to

variations in the specific welding parameters and conditions rather than

the calculated energy input. What is to be noticed from these micrographs

is the increasingly continuous ferrite network with increasing energy input.

Since ferrite may become brittle at low temperatures or will transform to

the more brittle sigma phase at high temperatures, a less continuous ferrite

network is desirable. This would suggest that a lower energy input would

be advantageous. Another observation from this series micrographs

is that a general increase in dendritic ferrite spacing occurs as heat

input is increased, effectively increasing the grain size of the weldment.

This increase in grain size is not advantageous either, since it results in

a somewhat weaker weld. This trend also supports the use of a lower heat

input

.

378
ocOU

roO

+-> • C +-> •r- 3

C.

<DC

4-> -rs- -t-> t- (T3 OJ <U

M- cn O C

•r“ ^ 1/1

-l-> fD 3 QJ o ss- u

C71

C

O)

CD 3

•r- CD 3 s= o

-Cl o

CO ro

CO CO JCC- U

03 -I— S- +-> cn-r- o s- S- -O o c:

•r- OJ E "o o s4-> OJ O 4-> JZ iC D_ •!-

QJ

S-cn coon

•oo

czI

CO

CNJ

CO

-Oo

C.'I r-^

•a

379

-III

Rod 3

46.1

KJ/in.
Streicher Corrosion Test

The Streicher (oxalic acid) test was performed as a preliminary corrosion

test to the more quantitative Huey test. Three representative micrographs

of the Streicher test results are shown in Figure 11. Figure llA shows

the step structure which is characteristic of the 316L austenitic stainless

steel base metal. The base metal is the only sample which formally passed

the Streicher test. Figure IIB shows a ditched structure partially

surrounding the primarily dendritic microstructure of this weld. This

micrograph also shows the large variation in dendrite widths which can

occur from bead to bead in the weldments. Figure IOC shows a heavily

ditched structure surrounding the cellular microstructure occurring in this

weldment. This series of micrographs also shows the tremendous variation

in microstructure that occurs in the same weldment. It is apparent that

the Streicher test, being a test to determine the susceptability of stainless

steel to sensitization (and thus to intergranular corrosion) cannot be

applied to weldments containing ferrite. Because of the severity of the

etch, it is difficult to determine whether the attack is associated with

a chromium depleted region adjacent to the grain boundaries or confined

to the ferrite phase itself.

Huey Corrosion Test

The Huey test determines the corrosion rate as a function of time.

The weight loss data was converted to weight loss per surface area of the

2 sample, and is plotted as weight loss vs. time for all weldments and base

2 metal on Figure 12 and 13. The t time dependence was found to give the

overall rate law and suggests an accelerated corrosion behavior in the

weldments.

380
3ol

heavily

ditched

cellular

structure
Corrosion

wt.

loss

(gms)/

surface

area

(cm)^

x

100

70

1

60 . V 12006 Rod 2

O — 12005 Rod I A 12009 Rod 5

12007 Rod 3

O — 12008 Rod 4

/

/

( Time ,hrs)^ x 10"^

2 Figure 12. Corrosion rate versus time for the Huey test data taken on five weldments, welded at a constant energy input but with

different electrodes.

382
Corrosion

[wt.

loss

(gms)/

surface

area

(cm)^J

x

100

Figure 13. Corrosion rate versus Time for the Huey test data taken on five weldments, welded with select electrodes and different

energy inputs. The 316-L base metal is also included on this

graph.

363
A general relationship between corrosion behavior and weldment ferrite

content can be seen in Figure 14 for the 240 hr. corrosion weight loss data

plotted as a function of % ferrite. The rate of corrosion rapidly increases

with /o ferrite up to a ferrite value of 6%, but then drops off to a relatively

constant rate independent of ferrite and comparable in corrosion rate to the

base metal itself. Since this data was taken from welds with various welding

parameters, a more representative plot of the data is shown on Figure 15

which shows the 240 hr. corrosion vs. % ferrite for the first five weldments,

welded with the same energy inputs. This figure also illustrates the same

type of behavior.

To explain this - corrosion behavior, the diagram shown in Figure 4 is

applied. Weight loss vs. chromium and nickel equivalent is plotted in

Figure 16 for the five different electrodes. A bar graph is also included

on Figure 17 showing the weight loss in 240 hours for each electrode. The

two points with low corrosion on the high chromium side of the dashed line

are from the high ferrite weldments while the other three points represent

the low ferrite containing weldments. Figure 16 shows that the corrosion

behavior changes drastically as the weldment composition crosses the liquidus

line. Thus the corrosion behavior is dependent on whether ferrite or

austenite is the proeutectic dendritic phase. Two forms of corrosion may be

responsible for the corrosion behavior of the stainless steel weldments:

intergranular corrosion and galvanic corrosion. Intergranular corrosion may

occur as the result of the formation of grain boundary or interdendritic

carbides which deplete the adjacent matrix of chromium. The chromium carbides

first form along the ferrite-austenite boundaries, then around the austenite

grain boundaries. Referring to Figure 6, which shows schematically the predicted microstructure for four specific electrode compositions, a material

with composition ’A’ would solidify having only austenite grain boundaries

384
I

00 CD

^

CJ o

o

00

CD

CM

o

0)o>o

a>>

<

<D

0>

001 ^ D8J0 aoD^jns /( slliB) sso| |m

]

9|Dy UOjSOjUOO

job

Figure

14.

Corrosion

rate

and

316-L

base

metal

vs.

%

Ferrite

for

all

ten

weldments.
C\J

O

00

CD

CM

O

s

0) "O

—I

QJ

c

a»eT3

(U3

01

cn >. > CJ£.

)-i 0) CLi u C

CO d>

S-I

O CO •H 4-1 CO CO O P P o

O oa CO

0) Vj3

&c

•H

001 X D8JD 80D|JnS /(SlIlB) SSO| \fA

]

8|Dy UOISOJJOQ

3ij6

%

Ferrite

(Average)
Cr 25 24 23 22 21 20

Ni 13 14 15 16 17 18

Figure 16. Fe-Cr-Ni ternary phase diagram. above which is plotted

t he corrosion rate of the first five weldments , welded

at a constant energy inpu t

.
240

Hrs.

Corrosion

Rote

wt.

loss

(gms)/surfoce

area

(cm)^

xlOO

2 3 4 5

Electrode

Figure 17. Bar graph showing the relative corrosion rates of the

five separate electrodes. Each electrode producing

the indicated amount of residual ferrite.

ooo JOO
available for carbide precipitation; for a composition 'B’ which passes

through the eutectic portion of the diagram, finely distributed ferrite exists

scattered throughout the matrix, producing a large area for carbide precipitation;

for composition *C’, ferrite is the primary phase to solidify giving additional

area for carbide precipitation; and finally for composition ’D’ there will exist

only a small amount of austenite-ferrite grain boundary surface area, since the

structure solidifies entirely as ferrite. Thus assuming that only intergranular

corrosion is responsible for the corrosion behavior, the corrosion rate would

increase from composition 'A' to 'B* to ’C* but then drop off for composition

'D*, explaining the corrosion behavior of the weldments welded with the five

separate electrodes as shown in Figure 16.

An alternate model can be suggested to explain the corrosion behavior

based on the use of the DeLong diagram shown in Figure 2. The line plotted

on the diagram represents the liquidus line of the portion of the phase diagram

of interest, Fe-Cr-Ni. On one side of the line, austenite is the primary

phase to solidify while on the other side, ferrite is the primary phase to

solidify. This line passes through a rectangular region outlining the composition limits of 316 stainless steel. The five points plotted on this diagram represent the compositions of the five electrodes used in this study.

Notice that the electrodes with high corrosion rates, electrodes 1,2, and 5

lie on the austenite side of the liquidus line while the electrodes with the

low corrosion rates, electrodes 3 and 4 lie on the ferrite side of the line.

Thus suggesting that liquating elements, sulfur, phosphorus, etc. , play a

role in the behavior as they do with the hot cracking behavior. Hot cracking

is more probable when austenite is the primary phase to solidify.

The second form of corrosion to consider, galvanic corrosion, may occur

when a potential difference exists between two dissimilar alloys or phases.

This potential difference causes electron flow between the two phases.

3b9
Usually the more corrosion resistant (more noble) material is protected at

the expense of the less corrosion resistant, more active material. The stainless steel of Figure 5a has almost no ferrite; there would be little galvanic

corrosion. The steel of Figure 5b has a higher ferrite content spread out

in the eutectic structure and being distributed as it is, the ferrite is high

in area and in close contact with the austenite, providing a means for galvanic

corrosion. Figure 5c has a higher cathode to anode area ratio providing an

even more accelerated attack. The sample in Figure 5d has a high ferrite

content, but the ferrite is not distributed as well as in the other microstructures, thus reducing the galvanic corrosive effect. In summary, corrosion

explanations could be used to explain the increase in corrosion rate with

increasing chromium equivalent (ferrite content) up to a ferrite content of

5 to 6% and then the corrosion rates drop off for higher ferrite levels as

experimentally determined in the Huey test.

A series of micrographs taken on a scanning electron microscope is shown

in Figure 18 showing the corrosion behavior as a function of residual ferrite

content. Figure 18a clearly indicates the intergranular corrosion of the

ferrite-free base metal. The weldment of Figure 18b contains 2.2% ferrite

(2.4 FN) and has a composition just on the high chromium side of the apex of

the three phase region. The weld shown in Figure 18c has 4.1% ferrite (4.5FN)

and has a composition at the high chromium end of the three phase region.

Notice the massive attack in some regions of the sample. The structure of

Figure 18d contains 101% ferrite (11.0 FN) and has a composition such that

ferrite is the only phase to solidify. This structure then, has a low

corrosion rate.

Figure 19 is a series of SEU micrographs showing the effects of energy

input on the weldment microstructure. These four weldments were all made

using electrode number 3; the energy input increases from Figure 19a to

Figure 19d. All of these welds have a high ferrite content which solidified

390
Q

TDC

fT3

C_) CDX O) O) OJ o +-> 4-> 4-> • 1

*r“ *1

—

C\J s- SS- C s_ —.

O) O) O) c M- 0l_ 4-

00 ^ c. O 00 o C\J 1— 1— -X Cl U CM E fO ro +-^ 00 +-> fO '—o c LD CNJ

•i- o 1—

00 -1— 03a "D

O oo 4-> o O

S- o <D oc cc

S- E O S- > u o O) 1—(

»—t u 00 1 1 >1 03 CTi oo

OJ oo CD

O o D D

IC

O

!-

O fU>

00 ^ oo

CL

C

ro •.- ^ 3

cn o O sz

S- 00 u

<C CD

O CD

O L-1

+-J OO O

jc c: D_ O

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O)

S-D

CD

Ll_

391

Rod 4

10.1%

ferrite
C) o

cn£Xo • O 4-> LD 13Q. dJ C

LU -M 00 O

QJ O) ^ Coc

cu -X

ra -M

t/1 <D

cn

c:

•1“CO

fO

<D

S_u

O

CO CO

CsJ

r3

CO

03 CU CO M C lO O

•rO

CO -F— o +-> S- -I- S- S- o *a u c:cu >iT3

cuO CD

ro

~ao

C'

ro

T3O

OC

OOoO

OCI I

M- CO O 03CU CO S- -C U

CL

C

03 T5-cn cu o -O

i~ +-> o

•r- cn E c O -r-

+-> 3 o o

JC ^ a. OO

CTi

03

CU

s-o

cn

392

Rod 3

49.9

KJ/in.

with

large

preheat
as high chromium-containing dendrites, leaving a chromium-depleted interdendritic liquid to solidify last. Notice how this chromium-depleted phase,

last to solidify, is selectively corroded leaving the dendritic solidification

pattern.

An important observation from these micrographs is that dendrite size is

sensitive to weld heat input, dendrite size increasing at higher energy inputs.

393
REFERENCES

1. C. D. Lundin and D. F. Spond, ’’The Nature and Morphology of

Fissures in Austenitic Stainless Steel Weld Metals,” Welding

Journal 55 , 356S (1976)

.

2. S. B. Melkuraov, "Influence of Specific Heat Input on the

Corrosion Resistance of the Metal of the Heat Affected Zone

if Ferritic-Austenitic Stainless Steel," Svar. Proiz, No. 4,

25 (1976) .

3. K. E. Pinnow and A. Moskowitz, "Corrosion Resistance of

Stainless Steel Weldments," Welding Journal , 49 , 278S (1970),

4. R. David Thomas, Jr., "Ferrite in Austenitic Stainless Weld

Metal," Arcos Corporation Report - March 3,’l978.

>J 5. W. T. DeLong, G. A. Ostrom, and E. R. Szamachowski , "Measurement and Calculation of Ferrite in Stainless Steel Weld

Metal," Welding Journal, 35 , 521S (1956).

6. C, J. Long and W. T. DeLong, "The Ferrite Content of Austenitic

Stainless Steel Weld Metal," VJelding Journal , 52 , 281S (1973).

7. G. M. Goodwin, N. C. Cole, and G. M. Slaughter, "A Study of

Ferrite Morphology in Austenitic Stainless Steel Weldments,”

Welding Journal , 51 (1972)

.

8. A. L. Schaeffler, "Constitution Diagram for Stainless Steel

Weld Metal," Metal Progress , 56 , 680 (1949) .

9. W, DeLong, G. Ostrom, and E. Szumachowski , "Measurement and

Calculation of Ferrite in Stainless Steel Weld Metal," Welding

Journal , 35, 526S (1956).

10. F. C. Hull, "Delta Ferrite and Martensite Formation in Stainless Steel," Welding Journal , 52 , 193S (1973).

11. R. Castro and J. J. De Cadenet, "Welding Metallurgy of Stainless and Heat-Resisting Steels," Cambridge University Press,

London (1975) .

12. D. Seferain, "Metallurge de la Foiidure," Dunod, Paris (1959).

394
REFERENCES (continued)

13. V. N. Zemzin and G. L. Petrov, "Influence of Ferrite Content

on the Properties of Austenitic Weld Metal," Svar. Proiz.,

, No. 5,6 (1967)

.

14. H. C. Ccimpbell, "The Ferrite Problem," Welding Journal, 54,

867 (1975)

.

15. H. F. Reid and W. T. DeLong , "Making Sense Out of Ferrite

Requirements in Welding Stainless Steels," Metal Progress, 84,

73 (1973) .

16. G. M. Goodwin, N. C. Cole, and G. M. Slaughter, "A Study of

Ferrite Morphology in Austenitic Stainless Steel V7eldments ,

Welding Journal , 51 , 425S (1972)

.

17. G. E. Linnert, "Welding Characteristics of Stainless Steels"

Source Book on Stainless Steels ; Section VI; Welding, p. 226,

American Society for Metals, Metals Park, Ohio (1976).

18. D. J. Kotecki, "Welding Parameter Effects on Open-Arc

Stainless Steel Weld Metal Ferrite," Welding Journal, 57,

109S (1978)

.

19. N. T. Williams, "Influence of Welding Cycles on the Ferrite

Content of Type 321 Austenitic Steel," British Welding Journal,

1^, 435 (1965)

.

20. R. A. Daemen, "High Recovery Electrodes for Welding Stainless

Steel," British Welding Journal , 14 , 550 (1967).

21. A. L. Schaeffler, "Selection of Austenitic Electrodes for

Welding Dissimilar Metals," Welding Journal , 26 , 601S (1947).

22. I. A. Zaks and R. S. Zhorova, "Effect of Various Fluxes on

the Solidification Cracking Resistance and Properties of

Austenitic Weld Metals," Svar. Proiz, , No. 7,32 (1974).

23. A. J. Zvanut and H. N. Farmer, Jr., "Self-Shielding Stainless

Steel Flux Cored Electrodes," Welding Journal , 51 , 775 (1972).

24. B. S. Payne, G. J. Dormer, and L. R. Haslip, "Lime Coated vs

Titanium Coated Stainless Steel Electrodes," Welding Journal,

41, 115 (1962)

.

25. N. C. Binkley, R. G. Berggren, and G. M. Goodwin, "Effects

of Slight Compositional Variations on Type E308 Electrode

Deposits," Welding Journal, 53 , 91S (1974).

395
REFERENCES (continued)

26. A Bernstein, J. Carlen, and L. Rick, "Influence of Phosphorus

and Sulfur on the Properties of the Weld Metal in Certain j Austenitic Stainless Steels," Welding Journal , 44 , 504S (1965).

27. C. V. Shirwaikar and G. P. Reddy, "Purging with Nitrogen in

the Welding of Austenitic Stainless Steels," Welding Journal,

54, 12S (1975) .

28. B. C. Odegard, "Some Observations of the Formation of Surface

Martensite in 304L Stainless Steel," Metallography, 7, 129 (1974)

29. T. Suzuki, H. Kojima, K. Sutaki, T. Hashimoto, and M. Ichihara,

"An Experimental Study of the Martensite Nucleation and Growth

in 18/8 Stainless Steel, Acta Met , 95 , 1151 (1977).

30. R. G. Berggren, N. C. Cole, G. M. Goodwin, J. O. Stiegler,

G. M. Slaughter, R. J. Gray, and R. T. King, "Structure and

Elevated Temperature Properties of Type 308 Stainless Steel

Weld Metal x^ith Varying Ferrite Contents," Welding Journal, P 167S (1978) . I 31. J, R. Hawthorne and B. H. Menke, "Influence of Delta Ferrite

Content and Welding Variables on Notch Toughness of Austenitic

Stainless Steel V7eldments," Naval Research Laboratory Report,

Washington, DC.

32. R. T. King, J. O. Stiegler, and G. M. Goodwin, "Relation i Between Mechanical Properties and Microstructure in CRE Type j 308 Weldments," Welding Journal , 53 , 307S (1974).

33. R. G. Thomas, "The Effect of Delta-Ferrite on the Creep Rupture

Properties of Austenitic Weld Metals," Welding Journal, 57,

232S (1978) .

' 34. Y. Arata, F. Matsuda, and S. Katayama , "Solidification Crack

Susceptability in Weld Metals of Fully Austenitic Stainless

Steels (Report II) —Effect of Ferrite, P, S, C, Si, and Mn on • | Ductility Properties of Solidification Brittleness," / \

Transactions Japanese Welding Research Institute , 6_, (1977) .

'

!

* I

’ 35. y. Arata, F. Matsuda, and S. Katayama, "Solidification Crack ? | Susceptability in Weld Metals of Fully Austenitic Stainless ^

'i Steels —Fundamental Investigation of Solidification Behavior

of Fully Austenitic and Duplex Microstructures and Effect of , Ferrite on Microsegregation," Transactions Japanese Welding

Research Institute, 5, 135 (1976) .

396
REFERENCES (continued)

36. J. A. Brooks and F. J. Lambert, Jr., "The Effects of Phosphorus

Sulfur and Ferrite Content of Weld Cracking of Type 309 Stainless

Steel," Welding Journal , 57 , 139S (1978).

37. Mrs. J. Honeycombe and T. G. Gooch, "Effect of Manganese on

Cracking and Corrosion Behavior of Fully Austenitic Stainless

Steel Weld Metals," Metal Const., and British Welding Journal,

4, 456 (1972)

.

38. D. M. Haddrill and R. G. Baker, "Microcracking in Austenitic

Weld Metal," BWRA Report B5/20/63 (1963).

39. M. O. Malone, "Sigma and 885F Embrittlement of Chromium-Nickel

Stainless Steel Weld Metals," Welding Journal , 46 , 241S (1967).

40. D. P. Edmonds, R. T. King, and G. M. Goodv/in, "Residual Elements

Have Significant Effects on the Elevated Temperature Properties

of Austenitic Stainless Steel Welds," ORNL Report TM-6177 (1978).

41. 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 Technology , 98 , 194 (1976)

.

42. L. A. James, "Crack Propagation Behavior in Type 304 Stainless

Steel Weldments at Elevated Temperatures," Welding Journal, 52,

173S (1973) .

43. P. Shahinian, H. H. Smith, and J. R. Hawthorne, "Fatigue Crack

Propogation in Stainless Steel VJeldments at High Temperatures,"

Welding Journal , 51 , 527S (1972)

.

44. R. T. King, D. A. Canonico and C. R. Brinkman; "Elevated Temperature Weldment Behavior as Related to Nuclear Design Criteria,"

Welding Journal , 54 , 265S (1975)

.

45. G. H. Rowe and J. R. Stewart, "Creep-Rupture Behavior of Type

316 Stainless Steel Weldments Prepared With and Without

Restraint," Welding Journal , 4

1

, (1962).

46. P. Shahinan, "Fatigue and Creep Crack Propagation in Stainless

Steel VJeld Metal," Welding Journal , 57 , 87S (1978) .

47. S. Ekerot, F. Bruk , and H. Fredriksson, "Pore Formation in the

Mechanism of Stainless Steel Weld Metals," Meta llographic Dept.

Royal Institute of Technology, Stockholm, Sweden.

48. N. C. Binkley, G. M. Goodwin, and D. G. Harman, "Effects of

Electrode Coverings on Elevated Temperature Properties of

Austenitic Stainless Steel Weld Metal," Welding Journal , 5

2

, 306S (1973) .

397
REFERENCES (continued)

49. C. D. Lundin, W. J. DeLong, and D. F. Spend, "Ferrite-Fissuring ® Relationship in Austenitic Stainless Steel Weld Metals,"

Welding Journal , 54 , 241S (1975) . .

50. J. C. Borland and R. N. Yonger, "Some Aspects of Cracking in

Welded Cr-Ni Austenitic Steels," British Welding Journal, 7,

22 (1960)

.

51. H. Frediks and L. J. Vander Toorn, "Hot Cracking in Austenitic 1 Stainless Steel Weld Deposits," British Welding Journal, 15,

178 (1968) .

52. W. T. DeLong, "Calibration Procedure for Instruments to Measure

the Delta Ferrite Content of Austenitic Stainless Steel Weld i Metal," Welding Journal , 52 , 69S (1973). i

53. R. A. Huseby, "Stress Relieving of Stainless Steels and the

Associated Metallurgy," Welding Journal , 37 , 304S (1958).

54. N. C. Code, G. M. Goodwin, and G. M. Slaughter, "Effect of Heat

Treatments on the Microstructure of Stainless Steel Weld Metal,"

Microstructural Science , 789 (1975)

.

55. E. M. Pronina et al, "The Effect of Strain Aging on the

Susceptibility of Austenitic Steel KH25N16G7AR (E1835) to

Cracking in the Heat Affected Zone of the Welded Joint," Suar.

Proiz. , No. 6,23 (1977).

56. W. E. White and I. LeMay, "Metal lographic Observation on the

Formation and Occurrence of Ferrite, Sigma Phase, and Carbides

in Austenitic Stainless Steel," Metallography , 51 (1970).

57. W. A. Baeslack, III, D. J. Duquette, and W. F. Savage, "Stress S Corrosion Cracking in Duplex Stainless Steel Weldments,"

Welding Journal , 57 , 175S (1978)

.

58. L. R. Poole, "An Unwanted Constituent in Stainless Weld Metal,"

Met. Prog. , 65 , 108 (1954) . . 59. O. H. Henry, M. A. Cordovi , and G. J. Fischer, "Sigma Phase in | Austenitic Stainless Steel Weldments," Welding Journal, 34, 75S

(1955)

.

60. W. A. Petersen, "Fine Grained V7eld Structures," Welding Journal, ' 52, 74S (1973) .

390
REFERENCES (continued)

61. R. J. Traman and D. Hardwick, "Some Effects of Heat Treatment

on the Rupture Properties of Three Stainless Steels," Iron and

Steel Institute Publication, £2 (1967)

.

62. H. F. Reid and W. T. DeLong , "Making Sense Out of Ferrite

Requirements in Welding Stainless Steels," Met. Progress, 103,

73 (1973)

.

63. R. D. Thomas, "Crack Sensitivity of Chromium-Nickel Stainless

Weld Metal," Metal Progress, _50, 474 (1946).

64. R. S. Brown and J. B. Koch, "Development of a Solidification

Crack Resistant Weld Filler Metal for a Stabilized, Fully

Austenitic Alloy," Welding Journal , 57 , 85S (1978)

.

65. D. M. Haddrill and R. G. Baker, "Effect of Cobalt on Susceptibility

of Welded Austenitic Steels to Heat Affected Zone Cracking During

Heat Treatment," B.W.R.A. Report, C104/1 (1964).

66. Mrs. J. Honeycombe and T. G. Gooch, "Effects of Microcracks on

Mechanical Properties of Austenitic Stainless Steel Weld-Metals,"

Metal Construction and British VJelding Journal , £, 140 (1973) .

67. J. Durnin and K. A. Ridal, "Determination of Retained Austenite

in Steel by X-Ray Diffraction," J. Iron and Steel Institute,

206, 60 (1968)

.

68. F. C. Hull, "Effects of Composition on Embrittlement of Austenitic

Stainless Steels," Welding Journal , 50 , 104S (1973).

69. C. L. Choi and D. C. Hill, "A Study of Microstructural Procession

in As-Deposited Weld Metal," Welding Journal , 57 , 232S (1978)

.

70. N. I. Kakhovski and A. M. Solokha, "A Cr-Ni Steel Containing

Nitrogen with a Low Carbon Content," Auto. Weld , 21 , 81 (1968)

.

71. J. K. Pokhodnya and A. P. Paltsevich,_ "Absorption of Nitrogen

During the Welding of Type 18Cr-9Ni Steel," Auto Weld , 23 , 9 (1970)

.

72. H. Astrom, B. Loberg, B. Bengtsson, and K. E. Easterling, "Hot

Cracking and Micro-Segrega tion in 13-10 Stainless Steel Welds,"

Metal Science , , 225 (1976)

.

73. M. C. T. Bystram and M. F. Stuchfiel.d, "Nickel Base Electrodes

for Plant Operating at Low and High Temperatures," British

Welding Journal , 11 , 427 (1964) .

399
REFERENCES (continued

)

74. E. H. Schmidt, "Low Temperature Impact of Annealed and

Sensitized 18-8," Metal Progress , 52 , 698 (1948).

75. J. R. Hawthorne and H. E. Watson, "Notch Toughness of

Austenitic Stainless Steel Weldments With Nuclear

Irradiation," Welding Journal , 52 , 255S (1973).

76. W. T. DeLong, "Ferrite in Austenitic Stainless Steel Weld

Metal," Welding Journal, 273S (1974).

77. R. Espy, "What Weld Metal for Type 316L Stainless Steel,"

Metal Prog. 109, (1964),

78. R. Gunia and G. Ratz, "The Measurement of Delta Ferrite in

Austenitic Stainless Steels," WRC Bulletin 132, August (1968).

79. I. Masumoto, K. Tamaki, and M. Katsuna, "Hot Cracking of

Austenitic Steel Weld Metal," J. Japan Weld. Soc . , £1, 1300 (1972)
Nonmetallics for Magnet Structures

401

NONMETALLICS FOR MAGNET SYSTEMS

National Bureau of Standards

403
1

I

i:
NONMETALLICS FOR MAGNET SYSTEMS

M. B. Kasen, J. G. Must and H. M. Ledbetter

Fracture and Deformation Division

and

Thermophysical Properties Division

National Bureau of Standards

Boulder, CO 80303

INTRODUCTION

NBS, in cooperation with the.U.S. laminating industry, has arranged for

commercial availability of special grades of NEMA type G-10 and G-11 glassepoxy industrial laminates intended for use in cryogenic magnet structures.

These materials, designated G-IOCR and G-llCR, differ from conventional G-10

and G-11 products only in that they are produced to a specification covering

all aspects of components and procedures, thereby minimizing product variability. The CR-grade laminates are commercial products meeting current NEMA

G-10 and G-11, MIL-P-18177C and Federal LP-509 specifications. G-IOCR is a

heat-activated, amine-catalyzed bisphenol A solid type epoxy resin laminate

reinforced with continuous-filament E-glass fabric, silane finished. G-llCR

os an aromatic-amine hardened bisphenol A liquid type epoxy resin laminate

reinforced in the same manner. The G-IOCR formulation was selected because

of its proven performance in cryogenic structures. The G-11 OR formulation

was selected to provide an alternative aromatic-amine cured laminate which

is expected to provide higher resistance to radiation damage.

This action was taken because variability in cryogenic properties among

conventional proprietary G-10 and G-11 laminates has caused concern as to

their suitability for demanding MFE applications in superconducting magnet

structures. Property variability problems arise because the room-temperature

electrical performance specifications, which conventional proprietary laminates

are required to meet, may be met by different laminate compositions and manufacturing processes. Thus, G-10 and G-11 laminates produced by different

408
companies can and do differ in the type of glass reinforcement, in the

weave of the glass fabric, in the epoxy system used and in other ways

which affect the cryogenic performance of the laminates. These problems

are eliminated by the CR grades.

Specifications for producing the CR grade laminates have been distributed to all U.S. laminating companies through the Industrial Laminates Subcommittee of the National Electrical Manufacturing Association,

Washington, D.C. Specifications will remain proprietary within NEMA.

Inquiries regarding details of the specifications should be addressed to

Mr. H. Colin Smith of NEMA.

The CR grades may be produced by any laminating firm wishing to

compete for the market. They are presently available in commercial

quanti ti es.

CR-grade laminates are produced to a natural beige color, as the

dye required to produce the fami 1 iar 'green color of proprietary G-10

and G-11 laminates is only cosmetic and has an unknown effect on radiation resistance. Standard laminated sheets range from 18 to 50 inches

in width and 24 to 120 inches in length. Standard thicknesses range

from 1/64 inch to 2 inches. Laminated tubes are available in a variety

of diameters and lengths. Both CR-grades are available in rod forms

produced by molding or by matching from square sheet stock.

Both G-IOCR and G-llCr are woven glass cloth composites. The warp

fibers are those which emerge from the loom, while the fill fibers are

406
those woven across them at right angles. The yarn count per inch is

43 ± 3 for warp and 32 ± 2 for fill. The normal direction of the composite is perpendicular to the cloth ply plane.

This report summarizes the results of cryogenic mechanical, thermal,

elastic and electrical property characterization of G-IOCR and G-llCR

laminates manufactured in the pilot plant of one industrial laminating

firm. A continuing study is examining the variability in properties

which might occur in normal plant- production and among the products produced by different firms to the same specifications.

The tolerance of G-IOCR and G-llCR to neutron and gamma radiation

at 4 K is being studied in an independent program at ONRL. Inquiries

should be directed to Mr. C. J. Long at that facility.

MECHANICAL PROPERTIES, THERMAL CONTRACTION, AND ELECTRICAL PROPERTIES

OF G-10 CR AND G-llCR

M. B. Kasen

MECHANICAL PROPERTIES

The effect of cryogenic temperatures on the mechanical properties

of G-IOCR and G-llCR laminates are summarized in Table 1 and are plotted

as a function of temperature on Fig. l(a-g). Cooling to cryogenic temperatures is clearly beneficial -- tensile and compressive strengths increase 70-100%, Young's moduli increase about 30%, ultimate tensile

failure strain is increased 140% and interlaminar (guillotine) shear

strengths increase 30-70%. In this test series, the G-IOCR product was

found to be mechanically superior to the G-llCR grade in all except

Young's modulus. There is also some indication that cooling below 76 K

may be deleterious to several of the G-llCR properties. These data do

not consider the factor of radiation resistance. If as expected, aminecatalyzed epoxy system of the G-llCR product suffers less degradation

407
INITIAL

MECHANICAL

PROPERTY

VALUES

OF

G-10

CR

AND

G-11

CR

(MANUFACTURER

A)

3 ZD - H

</) - c

O) c; c I—

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(W)

=

Warp

direction

of

reinforcement

(N)

=

Normal

to

reinforcement

*

=

Calculated

values
under irradiation than does the G-IOCR formulation, the performance

ranking of the two products might substantially change.

In common with conventional G-10 and G-11 products, the G-IOCR and

G-llCR laminates are reinforced with an unbalanced glass fabric, the

glass content in the fill direction of the fabric being about 75% of

that in the warp direction. The tensile and compressive properties in

the fill direction should, to a first approximation, be 75% of those

reported on Table 1 for the war.p direction. This remains to be verified.

Interlaminar shear properties (resistance to shear between layers

of the laminates) are important in superconducting magnet designs in

which the windings nest in slots machined into the laminates or in other

designs in which the electromagnetic forces place the laminates in

shear. The relevant data appear in Table 1 under the heading "Shear

Strength (Guillotine)." Data designated "Shear Strength (Short Beam)"

is not a relevant property for this application, as this test method

does not provide a true measure of interlaminar shear.

Interlaminar shear strength is a difficult property to measure accurately because the defined test conditions -- pure shear stress without

tensile or compressive stresses on the shear plane -- can only be approximated. The discrepancy noted on Fig. 1(g) between "unsupported"

and "side-supported" test methods indicates the possible error that can

occur even though the tests comply with an existing ASTM method which is

currently under revision. Data appearing on Table 1 and identified as

"side-supported" on Fig. 1(g) reflect improved methods developed at NBSBoulder and are believed to reflect the most accurate available data.

409
(a) Temperature dependence of ultimate (b) Temperature dependence of Young's

tensile strength, warp direction. modulus in tension, warp direction.

(c) Temperature dependence of Poisson's (d) Temperature

ratio in tension, warp direction. tensile strain.

dependence of ultimate

warp direction.

410

YOUNG'S

MODULUS,

E*

(psi)
csz

LU

crh*

C/5

00

00

cc

a.o

CJ

o

z

QC

C/5

00

00

QC

Q.O

o

(e) Temperature dependence of ultimate

compressive strength normal to the

reinforcement.

(T) Temperature dependence

compressive strength, warp

of ultimate

di rection

.

cs

cc<

TEMPERATURE, T(K)

(g) Temperature dependence of interlaminar shear (guillotine), warp

direction. Unsupported data are

invalid. See text.

Fig. 1. Temperature-dependence

of mechanical properties of GlOCR and G-llCR high-pressure,

glass-epoxy industrial laminates. Initial data, supplier

A, three specimens per data

point. Warp/fill ratio = NBS-Boulder data.
The "short beam" shear strength data appearing on Table 1 are

primarily indicative of overall laminate quality. The data indicate

that the tested laminates were equal in quality to the best commercial

products. Interlaminar shear strength values obtained with this test

cannot be equated with values obtained by the guillotine method because

the three-point flexure method of testing subjects the specimen to

tensile and compression stresses as well as shear with the state of

stress continually changing throughout the test. These data were obtained as part of a series of test evaluating the effect of changing

from a methacrylate chromium chloride (Volan) glass finish to a silane

finish as required by recent EPA directives. The results, summarized on

Fig. 2 indicate that the change of finish on the glass has substantially

improved the quality of both CR-type laminates.

The ability of the laminates to withstand severe thermal shock in

thick sections is of interest since cracking or delamination caused by

such shock could adversely affect the performance of the laminates in

magnet structures. The ability of G-IOCR and G-llCR to withstand

thermal shock was evaluated by repeatedly cycling one-inch cubes of each

material between 76 K and 390 K, after which degradation was assessed by

comparing the short-beam shear strength of specimens taken from the surface layers of cycled specimens with that of similar specimens taken

from uncycled cubes. The results, summarized on Table 2, suggest that

the CR-grade laminates are highly resistant to thermal shock.

The FY-79 nonmetal lies mechanical test program will complete the

characterization of the G-llCR and C-IOCR products in the fill direction.

A set of laminates received from a second manufacturer will be evaluated

to assess the variability in mechanical properties which might exist

412
Fig. 2. Comparison of short-beam shear strengths of G-IOCR

and G-llCR laminates produced with glass fiber having volan

(methacrylate chromium chloride) and silane finishes.

413
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cn OJ 0)

u

cs: > >

<c <t 01—1 •K

414
between laminates produced to the same specification at different

plants. Additionally, the mechanical properties of laminates produced

to the same specification but containing boron-free E-glass reinforcement will be screened. Samples of these laminates will be provided to

ORNL for 4 K radiation testing to determine if increased radiation

resistance due to reduction in boron content is sufficient to justify

additional development of this variant.

THERMAL CONTRACTION

The available data on contraction of G-llCR and G-IOCR laminates

are summarized on Table 3. These data, independently provided to NBS by

the Los Alamos Scientific Laboratory (R. I. Schermer), reflect total

thermal contraction from 295 K to 75 K.

A more complete characterization of thermal contraction is scheduled during FY-79.

Table 3.

TOTAL THERMAL CONTRACTION 295 K TO 75 K OF G-IOCR

AND G-llCR HIGH-PRESSURE, GLASS-EPOXY INDUSTRIAL

LAMINATES. (1

)

AL/L, mm/meter

Pi recti on G-1 OCR G-1 1 CR

Warp 2.35 ± 0.00^^^ 2.19 ± 0.07-

Fill 2.69 ± 0.03 2.43 ± 0.03

Normal 7.30 ± 0.15 6.20 ± 0.18

(1) Private communication, Los Alamos Scientific

Laboratory, 14 April 1978.

(2) Standard deviation of the mean of the measurements .

ELECTRICAL PROPERTIES

Available data on the effect of cryogenic temperatures on the volume

electrical resistivities of G-1 OCR and G-1 ICR laminates are summarized

41b
on Table 4. Data on the effect of cryogenic temperatures on the electrical strengths are summarized on Table 5. These data were provided 1

the Westinghouse Research and Development Center (R. N. Sampson) under

contract from NBS.

A more complete report on this work is being prepared by Westinghouse.

Table 4.

VOLUME RESISTIVITIES

AT LOW AND

OF G-10 AND G-11 LAMINATES

ROOM TEMPERATURES

Sampl

e

Thickness, Mils

Volume resistivity,

298K 77K

ohm-cm

4K

G-10 14.4 8.9X10^^ 4.0X10^^

G-10 20

14 17

9.3X10 1.5X10 4.1X10^^

G-11 12

1 5 7.3X10'^ 2.0X10^^

G-11 20 1.3X10^^ 6.1X10^^ 2.0X10^^

Table 5.

ELECTRIC STRENGTHS OF G-10 AND G-11 LAMINATES

AT LOW AND ROOM TEMPERATURES

Measurements with 60

0.5 in. and 1

Hz voltage at 0.5 kV/sec.

.5 in. diam. electrodes

rise.

Sample Thickness, Mils

Breakdown voltage stress, kV/mil at

298K 63K 4K

G-10 11 .87, 1.66,1.60

Two >1.52 1 .63 av.

G-10 14 1.16 1.26,1.19

1.23 av.

G-11 12 1.23

G-11 12.5 1.18

11.7 1 .25

1 .22 av.

416
Thermal Conductivity of G-ll)CR and G-llCR

Industrial Laminates at Low Temperatures

by

J. G. Hust

ABSTRACT

The thermal conductivity of G-IOCR and G-llCR is reported for temperatures from 2 to 300K. These materials are produced by a standardized manu- facturing specification designated by the suffix CR. Measurements were conducted in two directions: (a) Normal to the glass fabric plane and (b)

in tlie plane of the fabric (warp direction). The measurements and analysis

were conducted by the thermal conductivity integral method with an uncertainty of about + 5%.
I. INTRODUCTION

As part of the continuing cryogenic materials program, NBS, in cooperation with the U. S. laminating industry, has arranged for commercial production of component-specified NEIIA/ASTM type G-10 and G-11 industrial

laminates for cryogenic insulation and structural applications. This action was taken because variability in cryogenic properties of conventionally produced proprietary G-10 and G-11 laminates has caused concern in the

demanding specifications of superconducting magnets for magnetic fusion

energy applications. The cryogenic variability in G-10 and G-11 laminates

occurs because the room temperature NEMA specifications can be met by different glass reinforcements, weaves of the fabric, types of epoxy, and cure agents, as well is the fabrication procedures. The component-specified

laminates investigated are designated as G-IOCR and G-llCR.

N’BS is performing a thorough low temperature electrical, thermal, and

mechanical properties characterization of G-IOCR and G-llCR. This report

describes the thermal conductivity measurements performed on these laminates in two directions (to determine the anisotropy) and temperatures from

2 to 300 K.

II. MATERIAL AND SPECIMEN CHARACTERIZATION

G-IOCR and G-llCR are composites of fiber glass fabric and epoxy. The

fabric is woven from type E glass threads with 43 threads per inch (17 per

cm) in tiie warp direction and 32 threads per inch (12 per cm) in the fill

direction. The thread is composed of 75 filaments each of 0.00018 inch

diameter (0.00046 cm). The density of E glass is 2.4 g/cm^. The resin

Y-2000, used in G-IOCR, has an uncured density of 1.10 + 0.01 g/cm^. The

resin, HMS-2002 used in G-llCR, has an uncured density of 1.04 + 0.01

g/cm^. The densities of the cured resins are not known.

418
The thermal conductivity specimens are 3/A inch cubes. The G-IOCR

specimen has a density of 1.904 g/cm^ and the G-llCR specimen 1.956

g/cm^ with an uncertainty of + 0.002 g/cm^. The finished G-IOCR laminate contains 33-35% resin and G-llCR 27-31% by weight.

Using the above measured densities and weight fractions of the resins,

a value of 1.25 + 0.01 g/cm^ is calculated for the cured resin density

for both resins. This corresponds to a fiber content of 53% for G-IOCR and

38% for G-llCR by volume. These calculations were performed using the

following equations:

^'c -rWFr

V

c VF

+ .gWF

r

where . .ind V are density and specific volume (V=l/o), WF and VF are weight

fraction and volume fraction, and tlie subscripts designate composite (c),

resin (r), and glass (g).

Table 1. Dimensions, volumes, weights and densities of the G-IOCR

and G-llCR specimens

Dimensions (cm)

Ma t e r i a 1 Norma 1 to fabric

Warp

direction

Fill or woof

direction

volun^

(cm ) weight

(g)

density^

(g/cm )

G-IOCR 1.910 1.890 1.919 6.93 13.188 1.904

G-1 ICR 1.915 1.913 1.910 7.00 13.684 1.956

Thermal conductivity of each of these specimens was measured in the normal

and warp directions.

419
III. EXPERIMENTAL PROCEDURE AND DATA ANALYSIS

The thermal conductivity of the test blocks is determined in a modified "fixed-point" apparatus. This apparatus has been used previously to

determine thermal conductivity values near the fixed-points of cryogenic

baths (liquid helium, 4K; liquid nitrogen, 76K; dry ice-alcohol mixture,

194K; and ice-water mixture, 273K) for relatively rapid measurement to an accuracy of 10% or better. Modifications, as described by Hust and

Arvidson [1], were made to perform closely spaced measurements at temperatures from 2 to 300K. The specimen chamber of this apparatus is shown in

Figure 1.

A measurement is performed as follows: the chamber is placed into a temperature controlled bath and the Lower block heater is energized to generate a fixed temperature difference, -T, between the two blocks with a constant lieater power, Q. From a previous calibration experiment, the

amount :)f heat, 0^-, flowing through the compression bolts under these

conditions is kno\'n. The heat flow through the specimen of cross-sectional

area, A, and length, <• , is thus calculated as Og = Cj - Qj-. If the

temperature gradient is sufficiently small, the thermal conductivity, k, can be calculated from.

This relationship yields the correct value of k at T = (T^ + T2)/2,

regardless of the magnitude of ^ T = T 2 - Tj if k is either independent

of temperature between T and T 2 or a linear function of temperature,

k = aT. In this experimental arrangement it is necessary to use large

temperature differences in order to obtain data from 2 to 300K, thus, Eq

(1) would not yield correct values of k.

To circumvent this discrepancy we first resort to the defining

equation for k(T) for one-dimensional heat flow.

Eq. 1

420
Figure Specimen chamber of compression "fixed-point" thermal

conductivity apparatus.

421
Qs = k(T)A dT/dX Eq. 2

w’iiere dT/dX is the temperature gradient at any position along the specimen

and ’k.(T) is thermal conductivity at that temperature. Integrating this

equation from T| to T 9 we obtain

/A k(T)dT Eq . 3

Assuming some integrable functional form df(T dT for k(T) we obtain

Os? /A = f(l 2

) - fd^) Eq. 4

For convenience the function k(T) is chosen tvj be an appropriate lineir

series containing a number of undetermined parameters,

n k(T) = ^ ai dfp(T)/dT Eq. 5

i = l

Thus, we obtain

n

" y^a^(f j_(T 2

> - fi(Ti)) Eq. 6

i = l Only the linear parameters are unknowns in this equation. Thus,

a least squares fit of several experimental runs can be performed to

determine the a^ and, consequently, k(T) can be calculated from Eq . (5).

It is necessary to have at least as many independent runs as there are parameters but considerably more are obtained in actual practice. The

latter method (Eqs.(5) and ( 6 )) is referred to as the thermal conductivity

integral method, while the former (Eq. (1)) as the finite difference

method. Both of these methods are used here to analyze the experimental

data.

422
Further details of the experimental setup, determination of Qj.,

thermocouple calibration, ind the operating procedure have been described

in a previous sponsor report by Must and Arvidson [1]. A calibration mea- surement on a standard reference material is also included in the report by

Hust and Arvidson [1]. The measurements on the standard reference material

(SRM 735) indicate an overall system accuracy of about + 3% up to lOOK and + 6% above lOOK,

I V . RESULTS AN D DISCUSSION

Tne direct experimental data t)n the tvo specimens in two directions

are available from the author. These data were analyzed bv both the difference method an<J the thermal conductivity integral method. The results of

these calculations are illustrated in Figures 2, 4, 6 and 8. The deviat Loiio if tht- lirect experimental data from the integral method calculations

are illu-itrated in Figures 3, 5, 7 and 9. the function chosen for k(T) in

the integral method is

The r-'-nlting values of a^^ for the four specimens are presented in Table

2. T ilvjlar values, as calcualted from eq. 7, are given in table 3.

It is noted that these four sets of thermal conductivities are the same to within about + 20%, Because of this, a least squares fit was also

performed for the entire set of data and the results are presented in

Figures 10 and 11, the parameters are given in Table 2 and calculated

values in table 3. Figure 12 illustrates the smoothed data for each of

the four specimens. This plot reveals that G-IOCR and G-llCR exhibit

essentially the same thermal conductivity, but the warp directions has a conductivity about 20% higher in both cases.

6 i Eq.7

i=l

423
Table 2. Values of the parameters, a^^, of equation 7 for the four sets

of data on G-IOCR and G-llCR and for all of the data taken as a single si^t (composite mean).

G-IOR GR- IICR Composite Mean

Normal Warp Normal Warp

1 -122251.92 -29802.969 -28757.138 17929.675 -77070.427

255195.80 80143.031 73864.094 -22689.905 171354.28

1 - 1. 5 9650. 69 -23413.583 -28560.501 54878.220 -94121.693

4 50450.437 1551.1046 6340. 1721 -25877.220 26281.670

5 1 CC• 404.43393 -841.94250 4781.6591 -3613.2451

6 473.39836 -38.341244 59.450845 -298.89669 204.51475

NOTE: All values have been multiplied by 10 .

424
,1

[ Table 3, Values of thermal conductivity for the normal and warp directions

I of G-IOCR and G-llCR as calculated from equation 7 and the parameters in

I table 2.

Thermal Conductivity (W m ^ K ^)xlO^

Temp

(K)

Normal Warp Normal Warp Composite :

2.2 28.7 40.5 31.5 38.6 35.0

2.4 33.3 45.1 35.2 42.6 39.5

2.6 37.6 49.5 38.8 46 .

6

43.8

2.8 41.6 53.8 42.2 50.5 47.9

3.0 45.4 57.8 45.5 54.2 51.7

3.5 53.9 67.4 53.3 63.3 60.5

4 61.2 76.1 60.5 71.7 68.3

5 73.5 91.4 73.1 86.7 81.5

6 83.6 104 83.9 99.5 92. 5 8 99.6 126 102 120 no

10 113 143 116 136 124

15 139 175 143 163 152

20 160 200 163 182 173

30 192 241 193 211 209

40 218 276 216 236 240

50 238 308 237 260 266

60 256 338 255 284 291

80 286 395 289 331 336

100 312 448 321 378 376

120 336 498 352 423 414

140 361 545 382 466 450

160 386 590 412 507 486

180 413 633 442 546 520

200 440 674 471 582 553

240 500 752 529 649 620

280 565 823 587 708 685

300 600 857 735 717

42b
Thermal

conductivity.

TEMPERATURE, °F

-455 -450 -445 -440 -420 -380 -340 -300 -260 -160 +80

- 0.5

- 0.4

- 0.3

0.2

0 . 1

0.08

0.06

0.05

0.04

0.03

00

ri- c:

Temperature, K

Figure 2. Thermal conductivity values form G-IOCR, warp direction (+

obtained by difference methods; obtained from integral

method)

426

THERMAL

CONDUCTIVITY
-455

-450

-445

-440

-420

-380

-340

-300

-260

-160

+80

OJ

$-

13 +J

fO

O)Q.

E

o»

sqo oqBD sqo

%

‘ A/001 ( A - A)

427

Figure

3.

Deviations

of

observed

thermal

conductivity

integrals

(Y)

from

least

squares

fit

of

data

for

G-IOCR

(warp

direction)
Thermal

conducicivity

,

+80

TEMPERATURE, °F

-455 -450 -445 -440 -420 -380 -340 - 300-260 -160

1 1 1 1 1 1 1 1 1 1 r

Figure 4. Thermal conductivity values for G-IOCR, normal to fabric

(+ obtained by difference methods; obtained from integral

method)

428

THERMAL

CONDUCTIVITY
-455

-450

-445

-440

-420

-380

-340

-

300-260

-160

+80

- 13 -

429

Figure

5.

Deviations

of

observed

tliermal

conductivity

integrals

(Y)

from

least

squares

fit

of

data

for

G—

lOCR

(normal

to

fabric)
Thermal

conductivity

-455 -450 -445 -440 —I

TEMPERATURE, F -420 -380 -340 -300 -260 -160 +80

T r T 1 1—

r

T

Figure 6. Thermal conductivity values for G-llCR, warp direction

-u-jt ^ difference methods; obtained from integral

0.4

0.3

0.2

0.1

0.08

0.06

0.05

0.04

-10.03

D3

r+C

method)

14

430

THERMAL

CONDUCTIVITY
TEMPERATURE,

-455

-450

-445

-440

-420

-380

-340

-300

-260

-160

+80

'!

A/OOT ( A - A)

431

I

igure

7.

Deviations

of

observed

thermal

conductivity

integrals

(Y)

from

least

squares

fit

of

data

for

G-llCR

(warp

direction).
Thermal

conductivity,

-455 —I

TEMPERATURE, °F

-450 -445 -440 -420 -380 -340 -300 -260 -160 +80 —I

1 1 1 1 1 1

1 I r03C

Figure 8. Thermal conductivity values for G-llCR, normal to fabric

(+ obtained by difference method; obtained from integral

method)

432

THERMAL

CONDUCTIVITY
-455

-450

-445

-440

-420

-380

-340

-300-260

-160

+80 IiI

00D>

XJU3 O

•TDC

U

OO

CJ

l—i CT? ctE fO U HD nSI o

XJ

>1 ^ •iH U-i

0)

W

w 0) X)o ct3

U-I o“ o X w XJ c X o CT3 'r-( 0) XJ 1^

CO E > O

OJ V-. Q <4-1

o^

0»

Vj3

00

•r^

(i-i

"1 -

433

'-IICR

(normal

to

fabric)
Thermal

conductivity

TEMPERATURE

-455 -450 -445 -440 -420 -380 -340 - 300-260 -160 +80

T

1 1 1

1 1 1 1 1 1 T

0.5

0.4

0.3

0.2

0.1

0.08

0.06

0.05

0.04

0.03

3

*

Figure 10. Thermal conductivity values for G-IOCR and G-llCr, both

normal and warp directions (+ obtained by difference methods; mean values obtained by integral method)

- ' 0

-

434

THERMAL

CONDUCTIVITY,

Btu

ft
-455

-450

-445

-440

-420

-380

-340

-300

-260

-160

+80

I

•r; -

435

Figure

11.

Deviations

of

observed

thermal

conductivity

integrals

(Y)

least

squares

fit

of

all

data

on

G-IOCR

and

G-OllCr,

both

normal

and

warp

directions.
I

i

43G

TEMPERATURE,

K

Figure

12.

Thermal

conductivity

from

2

to

300 K

of

G-IOCR

and

G-llCR

in

the

warp

and

normal

directions.
Dynamic Elastic Modulus and Internal Friction in G-IOCR and G-llCR

Fiberglass-Cloth-Epoxy Composi tes*

H. M. Ledbetter

Fracture and Deformation Division, MBS, Boulder, Colorado, USA

ABSTRACT

Young's moduli were determined dynamically for two fiberglasscloth-epoxy composites in the warp, fill, and normal directions between

room temperature and liquid-nitrogen temperature. Dynamic internal

friction relates inversely to dynamic modulus in the studied materials.

The experimental arrangement consisted of a Marx three-component oscillator at frequencies between 40 and 90 kHz.

Key Words: Composites; elastic constants; glass-epoxy; internal friction

Young's modulus.

*Contribution of NBS, not subject to copyright.
INTRODUCTION

Solid-state elastic constants fill many needs. Engineering design

calculations require them for estimating load-deflection and thermoelastic stress. Derived from fundamental interatomic forces, elastic

constants index both cohesion and strength. They relate to other

physical properties such as specific heat and thermal expansion, all of

which help define a solid's equation of state.

A composite's elastic constants serve further purposes. Their

accurate measurement permits tests of theories relating composite properties to constituent properties. Their tensor character permits detecting fiber-induced mechanical anisotropy. Elastic constants may also

1 2 facilitate quality control and incipient-failure detection. For some

composites, empirical relationships exist between elastic constants and

3 important mechanical properties.

A composite's internal friction, which is a companion property to

the elastic constants, has several possible sources:

1. fiber properties

2. matrix properties

a. intrinsic (for example, "molecular engineering" in the case of

resins)

b. Extrinsic (for example, fillers)

3. fiber-matrix interface

a. relative displacement

b. local ized stress

4. cracks and voids

Besides depending on direction, internal friction usually varies with

both deformation mode (flexure, torsion, etc.) and frequency. The

433
present study used an extensional, or Young's-modulus-type, deformation,

which has alternating uniaxial tensile and compressive regions with

corresponding transverse Poisson contractions or expansions. Internal

friction not only characterizes composites, but it may serve to detect

2 incipient failure. In engineering design, higher internal friction

improves dimensional stability by minimizing stress-induced deflections.

The presently reported study had two principal purposes. First,

apply to composites an experimental technique used previously almost

4 exclusively for noncomposites for determining dynamic Young's modulus

and internal friction. Second, determine for G-IOCR and G-llCR composites their room- temperature Young's moduli and the change of these

moduli with decreasing temperature.

The experimental technique, using a three-component oscillator,

offers several possible advantages over static methods: relatively

small, simple-geometry specimens; simple and relatively small experimental arrangement, reducing refrigeration costs during cooling experiments; readily obtained one-percent inaccuracies; elastic-constant

changes due to temperature, pressure, magnetic field, etc., easily

5 measured to a few parts in 10 ; nearly simultaneous i nternal -friction

measurements. However, the technique does not yield accurate dilatationtype elastic constants, such as Poisson's ratio and the bulk modulus,

compared with directly measured shear-type elastic constants such as

Young's modulus and the torsional modulus.

EXPERIMENTAL

Materials

Table 1 summarizes studied-material properties. Materials were

received as green plates 2.5 cm (1 in) by 20.3 cm (8 in) by 20.3 cm

from manufacturer A.

439
Specimens

The experiments used circular rods, 0.40 and 0.48 cm diameter.

Lengths varied from 2 to 5 cm depending on material, fiber orientation,

and chosen resonance frequency. Three specimens were studied for each

material and orientation.

Elastic-Constant Measurements

Determining Young's modulus, E, required measuring the sound velocity, V, and the mass density, p, and using the relationship

c 2 .2,2

E = pv = pf^X (1)

where f^ denotes specimen resonance frequency and X denotes resonance

wavelength. At half-wave resonance, X = 21^ (where 1^ denotes specimen

length), and

E = 4pf^l2 (2)

5 6

A Marx ’ oscillator (shown schematically in Fig. 1) yielded the experimental resonance frequency according to

m e2f 0 0 (3)

where m denotes mass, o denotes the complete oscillator, and i denotes

oscillator components: quartz-crystal driver, quartz-crystal gage, and

specimen. It follows simply from Eq. (3) that

f^ = f^ + (f^ - f^) m /m

s 0 ' 0 q' q' s (4)

440
where subscript q denotes quartz crystals. The quartz-crystal length

fixes fq. A maximum voltage V^, corresponding to a maximum strain

amplitude at the gage-crystal midpoint, detects Sweeping frequency

f reveals the voltage maximum. Extensional elastic waves were launched

q along the rod axis by using -18.5° X-cut quartz crystals. Multiplying

the right side of Eq. (2) by the well-known Pxayleigh factor

R

= 1 +

n7Tvd\2

)

(5)

introduced the necessary short-rod correction, where n =

21 ^/X, v

denotes Poisson's ratio, and d denotes either the diameter of circular

rods or d = j

+ b ) for axb rectangular rods. Detecting resonance

at the first and third harmonics and extrapolating to zero yielded the

zero-frequency Young's moduli.

Internal -Fri ction Measurements

For a sharp, Lorentzian resonance peak (as shown schematically in

Fig. 2) the composite oscillator's intenal friction, is

^0 - (^2 - ( 6 )

the resonance-peak width at 1//? of maximum in a strain (proportional to

Vg) versus-frequency plot. Equation (6) holds for cf) = Q”^ up to 0.1 or

0.2, where 4> relates the complex, real, and imaginary elastic moduli as

shown in Fig. 3. Experimentally one observes M*, the modulus in phase

with strain. But one usually seeks j The usually more-rel evant

storage modulus. For zero internal friction, = M*.

Similar to Eq. (3)

,

m Q

-1

= X^m.Q. 1^1

1

1 (7)

441
Thus, a relationship similar to Eq. (4) yields the specimen's internal

friction

Q"^ = Q"^ +

(Q"^ - Q"b m /m (8) ^s 0 0 q

^ q s

^ '

RESULTS

Table 2 contains room- temperature and liquid-nitrogen-temperature

internal-friction values and Young's moduli at the quartz-crystal

resonant frequency (40 to 90 kHz). Table 3 and Figs. 4 and 5 show

Young's modulus versus temperature for the two composites.

DISCUSSION

Several observations follow from the results of this study:

1. Young's modulus can be measured more accurately dynamically

than statically, as shown by the average dynamic-case uncertainty of 3

percent.

2. Young ' s-modul us anisotropy, E(max)/E(min) , was about two, a

surprisingly low value. In-plane anisotropy was even lower.

3. Internal -friction anisotropy was about three. Internal friction was always higher in the normal direction, which tends to sample

the matrix. In-plane internal friction was higher in the fill direction

than in the warp direction.

4. Higher Young's modulus alv/ays corresponds to lower internal

friction, and vice versa.

5. All Young's moduli show regular behavior during cooling to

76 K: linearity at higher temperatures, steady increase, continuously

decreasing slope, tending toward zero slope at zero temperature.

6. During cooling. Young's modulus changes most in the normal

direction and least in the warp direction.

442
7. Internal friction decreased with low temperatures to a value at

76 K about 10 to 20 percent of its room- temperature value.

ACKNOWLEDGMENT

This study was supported in part by the U.S. Department of Energy.

R. E. Schramm and M. B. Kasen furnished materials for study and provided

unpublished data. M. Austin and G. Maerz made the measurements.

REFERENCES

1. K. Thomas, D. E. Meyer, E. C. Fleet, and M. Abrams, J. Phys. D. (Appl. Phys.) 6, 1336-1352 (1973).

2. L. E. Nielsen, J. Compos. Mater. 9^, 149-1 56 (1975).

3. A. Vary and K. J. Bowles, NASA Tech. Memo. TMX-73646 (1977).

4. W. H. Robinson, S. H. Carpenter, and J. L. Tallon, J. Appl.

Phys. 1975-1981 (1974). References 4-12 therein.

5. J. Marx, Rev. Sci. Instrum. 2^, 503-509 (1951).

6. W. H. Robinson and A. Edgar, IEEE Trans. Son. Ultrason. SU-21 , 98105 (1974).

7. A. S. Nowick and B. S. Berry, Anelastic Relaxation in Crystalline

Sol ids

,

Academic, New York (1972), p. 18.

443
Table 1. Properties of studied materials.

Mass Density Fiber Volume

Material Fiber Matrix (g/cm^) Fraction

G-IOCR 7628 E-gl ass cl oth bisphenol epoxy 1.90 0.64-0.68

G-llCR 7628 E-glass cl oth bisphenol A, aromatic amine 1.95 0.64-0.68

• r

Table 2. Dynamic Young's moduli, E, and internal friction, Q~\

of G-IOCR and G-llCR.

Temp

(K)

Di recti on G-IOCR

E(GPa) 0

1

O' G-llCRE(GPa)

1 01

O'

295 warp 29.5±1.0 68.3±6.6 31.4±0.4 147.0±34.9

fill 26.8±1 .0 114.0±7.8 27.7 161 .3

normal 14.0 228.6 15.6±0.6 406.5

76 warp 34.5 14.6 37.6 33.4

fill 33.3 17.3 -- -- normal 21.8 19.6 23.4 --

76/295 warp 1.17 0.21 1.20 0.23

dimensionless

ratios fill 1.24 0.15 -- -- normal 1.56 0.09 1.50 --

444
Table 3. Dynamic Young's moduli of G-IOCR and G-llCR at selected

temperatures in units of GPa.

Temp warp

G-IOCR

fill normal warp

G-llCR

normal

295 29.5 26.8 14.0 31.4 15.6

290 28.9 27.3 14.5 31.9 15.9

280 29.2 27.8 15.0 32.3 16.3

270 29.5 28.4 15.5 32.7 17.1

260 -- 28.9 16.0 33.1 17.6

250 30.2 29.4 16.5 33.4 17.8

240 30.5 29.8 17.0 33.7 18.5

230 30.8 30.2 - 17.4 34.0 19.0

220 31.1 30.6 17.9 34.3 19.4

210 31.3 30.8 18.1 34.5 19.6

200 31 .8 31 .5 18.8 34.9 20.1

190 31.0 31.7 19.2 35.1 20.5

180 32.3 32.0 19.7 35.2 20.8

170 32.6 32.2 19.9 35.5 21.1

160 32.9 32.4 20.2 35.8 21.4

150 33.1 32.6 20.5 36.0 21.5

140 33.2 32.7 20.6 36.1 21.6

130 33.5 32.9 21.0 36.4 21.9

120 33.6 32.9 21.2 36.5 22.1

no 33.7 33.0 21.3 36.7 22.4

100 33.9 33.0 21.5 36.9 22.3

90 34.0 33.0 21.6 37.0 22.6

80 34.3 33.1 21.7 37.2 22.7

76 34.5 33.3 21.8 37.6 22.8

445
Figure Captions

Fig. 1. Marx composite oscillator for determining standing-wave

(resonant) sound velocities.

Fig. 2. Forced-vibration Lorentzian-shaped resonance peak. Adapted

from Nowick and Berry. 7

Fig. 3. Complex-plane diagram showing vector relationships among

stress, strain, M* = complex modulus, = Storage

modulus, = ^imaginary

^ modulus, and <p = loss angle

internal friction. Adapted from Nowick and Berry.

Fig. 4. Temperature variation of three principal Young's moduli of G

lOCR glass-cloth-epoxy composite.

Fig. 5. Temperature variation of two principal Young's moduli of

glass-cloth-epoxy composite.

446
Marx composite oscillator for determininn standinn-wave

(resonant) sound velocities.

Fig. 1

.
Strain

Fig. 2. Forced-vi brati on Lorentzian-shaped resonance peak.

Adapted from Nowick and Berry. 7

448
Fig. 3. Complex-plane diagram showing vector relationships amona

stress, strain, = complex modulus, = storage modulus, flp = ^'-jniag-jnapy

^ ^°ss modulus, and <h = loss

annle = internal friction. Adapted from ilowick and Berry.

449
E(T)

E(295)

TEMPERATUKE (K)

Fig. A. Temperature variation of three principal Young's moduli

of G-IDCR glass-cloth-epoxy composite.

4j0
E(T)

E(295)

Fig. 5. Temperature variation of tv;o principal Young's moduli

of G-llCR glass-cloth-epoxy composite.

451

Technology Transfer
i

'i&i

•j
STRUCTURAL ALLOYS FOR SUPERCONDUCTING

MAGNETS IN FUSION ENERGY SYSTEMS

Nati onal -Bureau of Standards

Structural Alloys for Superconducting Magnets in Fusion Energy Systems*

H. I. McHenry and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

*Publication of NBS, not subject to copyright.

Submitted for publication in Nuclear Engineering and Design.
ABSTRACT

The behavior of selected alloys for superconducting magnet structures in fusion energy systems is reviewed with emphasis on the following

austenitic stainless steels (AISI grades 304, 310S and 316), nitrogenstrengthened austenitic stainless steels (types 304LN, 316LN and

21Cr-6Ni-9Mn) and aluminum alloys (grades 5083, 6061 and 2219). The

mechanical and physical properties of the selected alloys at 4K are re- viewed. Welding, the properties of weldments and other fabrication con- siderations are briefly discussed. The available information suggests

that several commercial alloys have adequate properties at 4K and sufficient fabrication characteristics for the large magnet structures needed

for fusion energy systems.

Key Words: Aluminum alloys, cryogenics, elastic properties, fabrication,

fatigue, fracture, fusion energy, stainless steel,

superconducting magnets, tensile properties, thermal

properties, welding.
1 Introduction

Alloys will be used for the main structural members of the superconducting

magnet systems needed for fusion devices. Designs of tokamak and mirror machines

that require superconducting magnets are currently in the conceptual stage

(except for MFTF), but many design features that strongly influence materials

selection are reasonably well established. The principal features governing

alloy selection for structural applications are the enormous size and stored

energy of the magnet systems, the extremely high forces exerted by the magnets,

the massive structural elements needed to restrain these forces, the limited

space available for the structure, and the need for accessability to install and

periodically remove the blanket and shield systems —systems that are completely

surrounded by magnets and the support structure. These features result in the

need for alloys with sufficient strength and stiffness to perform the required

structural functions within the space available, sufficient fatigue and fracture

resistance to operate safely, and sufficient f abricability to permit manufacture

and assembly of the components.

The choice of materials is further restricted by the adverse operating en- vironment of the magnet system. The alloys are subjected to temperatures down to

4K and must not be susceptible to low temperature embrittlement. The presence of

high magnetic fields of controlled shape cautions against the use of ferromagnetic materials in certain applications because they would be subjected to higher

forces due to the magnetic field and would result in field distortion. The rapid

magnetic field changes associated with cyclic operation of the poloidal field

colls cause heat inputs to the structure from eddy current losses. Consequently,

alloys with high electrical resistivity are usually preferred. The radiation en- vironment eliminates use of alloying elements, such as cobalt, which transmute to

long half life products. With this exbeption, radiation does not appear to be a significant factor in the selection of structural alloys because dosage levels

sufficient to degrade structural performance probably cause unacceptable deterioration of other system materials. Finally, the microstructure of the material

must be stable when exposed to the operating strains, temperature, and radiation;

e.g. the martensite transformation in metastable austenitic stainless steels produces a brittle, ferromagnetic phase and the attendant volume expansion causes localized areas of stress concentration.

459
The constraints on materials selection caused by size requirements,

space limitations and operating environment significantly reduce the number of candidate alloys. Further limitations arise due to considerations

related to experience, cost and availability. Currently, the leading

candidates are the austenitic stainless steels and aluminum alloys. In

this paper, the properties of the alloys considered most suitable for

liquid helium applications are reviewed with emphasis on their properties

at 4K and on their fabrication characteristics relevant to fusion magnet

systems.

This review draws heavily on research conducted at the National

Bureau of Standards; particularly a three-year program on materials for

superconducting machinery sponsored by the Defense Advanced Research

Projects Agency (1), a survey of low temperature materials for magnetic

fusion energy (2) and a continuing program on materials for fusion energy

magnets (3) sponsored by the Office of Fusion Energy of the U. S.

Department of Energy. For background relating to the properties

discussed the reader is referred to reviews by Fickett (4) on the

physical basis for each property and by McHenry (5) on fracture mechanics

and its application to cryogenic structures.

2. Annealed Austenitic Stainless Steels

Austenitic chromium-nickel stainless steels, including the nitrogen

strengthened grades discussed in more detail in Section 3, are the most

widely proposed alloys for structural applications in superconducting

magnet systems. These steels have the best combination of strength,

stiffness and toughness at 4K. The physical properties also offer advantages over competing materials. Specifically, the modulus is high, the

thermal expansion is close to that of the copper-stablized conductor,

magnetic permeability is low and the electrical and thermal conductivities are low. The principal disadvantages are the high cost of construction and problems arising from mlcrostructural stability and control.

The AISI 300-series stainless steels, particularly types 304 and

304L, are widely used in cryogenic applications. The designations and

460
compositions of commonly used grades are summarized in Table 1. These

grades have moderate strength, excellent toughness, good fabrication

characteristics and are available in a variety of product forms. The

ASTM specifications for the various product forms are listed in Table 2.

Several 300-series grades not listed in Table 1 are covered by ASTM

specifications and can meet ASME Pressure Vesels Code requirements for

cryogenic vessels. These grades are not described in this chapter be- cause they offer no particular advantage for superconducting magnet

structures and consequently their service experience and property data at

4K are limited.

The alloys listed in Table 1 can be compared with respect to AISI304, the basic 19Cr-9Nl stainless steel. AISI 316 contains 2-3% Mo and

slightly higher nickel and thus has greater austenite stability than AISI

304. AISI 304L and 316L are low carbon modifications of 304 and 316

respectively; low carbon is desirable to avoid sensitization, the grain

boundry precipitation of Cr-carbides. Type 310S is a 25Cr-20Ni alloy

that has lower carbon (.08%C) than AISI 310 (.25%C). The high alloy

content provides austenite stability and greater strength and thus 310S

is useful where dimensional stability, non-magnetic behavior and higher

strength are desireable.

The low temperature properties of the 300-series austentitic stainless steels have been compiled in cryogenic handbooks (6,7), in a review

article by Brickner and Defilippi (8) and in suppliers publications

(9,10). In this section, the general trends in strength, toughness and

fatigue resistance are summarized as a function of temperature. Care

should be taken in using the data presented here or in data compilations

because there are significant heat-to-heat variations in the mechanical

properties of the 300-series stainless steels, particularly in the yield

strength and toughness. The principal causes of these variations are interstitial content, i.e. carbon and nitrogen, and mechanical deformation history, i.e. resulting grain size. The influence of nitrogen within the limits of 0 to 0.08 weight percent is particularly significant

with the higher nitrogen contents causing substantial increases in yield

strength.

4C1
2.1 Tensile Properties

The yield and ultimate strengths of AISI grades 304, 310 and 316 are compared

in Figure 1 for temperatures ranging from 4 to 300K. Notice that type 304 has

the lowest yield strength and the highest ultimate strength at all temperatures. This is due in part to the martensite transformation which occurs more readily in 304 than in 316 or 310. Since the transformation is strain induced,

the yield strength is not significantly Influenced by martensite, but the ultimate strength is increased. '

The yield and ultimate strengths of AISI grades 304 and 304L are compared

in Figure 2. Type 304 exhibits a slightly greater increase in strength than

type 304L as temperature is reduced. The loss of strength associated with re- duced carbon content is considered to be characteristic of all the stainless

steels, i.e. 304L vs 304, 316L vs 316 and 310S vs 310, and is attributed to the

strengthening effect of carbon.

The ductility of the annealed 300-series stainless steels is generally ex- cellent at cryogenic temperatures. Elongation and reduction-of-area tend to

drop with decreasing temperatures, but values generally exceed 30%.

2.2 Toughness

The austenitic stainless steels retain excellent toughness at cryogenic

temperatures. The Charpy V-notch impact toughness does decrease with temperature but in all cases known to the authors the toughness far exceeds the 0.38mm

lateral expansion requirement of the ASME Boiler and Pressure Vessel Code (11)

and a ductlle-to-brlttle transition is not exhibited. The minimum temperature

for Charpy testing is about 20K, becaure adiabatic heating causes specimen

temperatures in excess of 20K even when the test temperature is 4K.

Slow-strain-rate tests using notched tensile specimens indicate that notch

sensitivity does not develop at 4K (6). The data summarized in Figure 3,

indicate that the notch-to-unnotched ratio of tensile strengths is approximately equal to 1 and the ratio of notched tensile strength to unnotched yield

strength generally exceeds 2.

462
Tobler (12) used J-integral methods to measure the fracture toughness of AISI grades 310 and 316 at temperatures to 4K. The toughness,

Kic(J), at 4K is 230 MPa /m for 310S and 460 MPa/ffT for 316, and in

both alloys the toughness at 4K exceeds the toughness at room temperature. The ratio of toughness to yield strength is sufficiently high to assure gross ductile deformation prior to fracture.

2.3 Fatigue

Strain cycling fatigue properties at 300, 76 and 4K have been measured for AISI grades 304L and 310 by Nachtigall (13) and for grades 304L

and 316 by Shepic and Schwartzberg (14). The results for 304L which were essentially the same in both investigations and for 316 are shown in

Figure 4. The fatigue resistance of 310 and 316 were superior to that of

304L, particularly in low cycle fatigue (less than 10^ cycles). For

each alloy, the fatigue resistance at low temperatures is superior to the

fatigue resistance at room temperature except at the highest strain

ranges where failure occurred in less than 1000 cycles.

The fatigue crack growth behavior of AISI grades 304, 304L, 310S and

316 has been determined at 295, 76 and 4K by Tobler and Reed (15). The

data for 304L are shown in Figure 5. Notice that the growth rates are essentially the same at 76K and 4K; similar behavior was observed in each

of the other alloys. Also, the growth rates at 76K and 4K are slower

than those at room temperature. The best-fit lines through the 304L data

at 295K and at 76K and 4K form the approximate scatter bands for the data

on 304, 310S and 316 at 295, 76 and 4K. 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.

463
2.4 Physical Properties

The physical properties of the 300-series stainless steels at 295,

76 and 4K are summarized in Table 3. The thermal properties have been

taken from the best-fit lines used in the LNG Materials Handbook (7), and

the original sources are given for the elastic properties (16), electrical resistivity (17) and magnetic permeability (18).

3. Annealed Nitrogen-Strengthened Austenitic Stainless Steels

Nitrogen has a pronounced effect on the yield strength of austenitic

stainless steels, particularly at cryogenic temperatures. For example,

the yield strength of 304N (AISI 304 with deliberate additions of nitrogen) is approximately three times greater than the yield strength of 304

at 76K, the ultimate strengths are approximately equal and the elongation

of the 304N is superior at temperatures below 220K (20). This remarkable

improvement in mechanical properties occurs in many nitrogen strengthened

grades. Consequently, significant interest is developing in these alloys

for future applications in superconducting magnet systems.

The designations and compositions of the commercially available

grades that offer promise for liquid helium applications are summarized

in Table 4. Several other nitrogen strengthened grades have been developed and are covered by ASTM specifications such as A240 and A412. However, the availability and property data on these grades are limited.

Several additional grades are commonly used in Europe where 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 and 316LN.

Two basic types of nitrogen strengthened stainless steels are suitable for cyrogenlc service: 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 the U.S. 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

464
provides austenite stability and can thus be used to replace part of the

Ni content. Manganese also Increases the solubility of nitrogen in

austenite and thus permits higher nitrogen levels-0.4% maximum nitrogen

is typical.

3.1 Tensile Properties

The yield and ultimate strengths of the Cr-Ni-N 300-series stainless steels are summarized in Figure 7. The solid lines and the broken

lines represent the data of Randak et al.(21) for 304LN and 316LN, res- pectively. The various symbols depict the data for seven other investigations (20,22-27), each on one of the following alloys: 304LN, 304N and

316LN. As is the case for 304 and 316, Figure 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. There is relatively little scatter in the yield strength data considering that data on 17 heats from a variety of sources are plotted. The nitrogen contents of the 17 heats

range from 0.09 to 0.17%.

Of the four nitrogen-strengthened Cr-Ni-Mn stainless steels listed

in Table 4, the 21Cr-6Ni-9Mn alloy has been most thoroughly evaluated

(28-32). The tensile and yield strength data are summarized in Figure 6.

This alloy typically has about 25% greater yield strength than the

Cr-Ni-N alloys. Representative tensile data on the other N-strengthened

Cr-Ni-Mn alloys that have been tested at 4K indicate that yield strength

increases and ductility decreases with decreasing temperature; and thus,

fracture resistance must be considered when selecting these alloys for

liquid helium service.

3.2 Toughness

The N-strengthened stainless steels have excellent toughness at room temperature. However, as shown in Figure 8, significant tougliness losses

generally occur as the temperature is reduced (25). Read and Reed (25)

observed significant variations in the toughness of a single piece of

46G
2 lCr-12Ni-5Mn alloy. As received mill-annealed material having an intergranular micro-constituent had a plane strain fracture toughness at 4K at

111 MPa/ral A high temperature (1177 C, 1.5 h) anneal dissolved the microconstituent and increased the toughness to 176 MPa/nT. Thus, certain

N-strengthened grades can be used at 4K, but care should be taken to assure that the alloy selected has satisfactory toughness at 4K for the

applicable melting practice, product form and heat treatment.

3.3 Fatigue

The strain cycling fatigue behavior of 21Cr-6Ni-9Mn has been measured at 295, 76 and 4K by Sheplc and Schwartzberg (14). The results At

4K, shown in Figure 4, indicate that the 21Cr-6Ni-9Mn alloy has fatigue

strength better than 304L but not as good as 316. Thus, the significantly

higher strength of the N-strengthened grade does not result in a comparable improvement in fatigue life.

The fatigue crack growth behavior of several N-strengthened grades

has been studied at cryogenic temperatures (25,28). The results for the

high nitrogen Cr-Ni-Mn alloys depend largely on the relative austenite

stability of the alloys. For the least stable alloy evaluated,

18Cr-3Ni-13Mn, the fatigue crack growth rates at 4K are 50 times faster

than those for the most stable alloy evaluated, 21Cr-12Ni-5Mn, and 10

times faster than the growth rates for 21Cr-6Nl-9Mn. For the

2 lCr-12Ni-5Mn alloy, the fatigue crack growth rates fall within the same scatter band as the 300 series (N<.08%) alloys at 4K. For the Cr-Ni

alloys, comparison of 304 and 304L with 304N and 304LN shows that the

nitrogen-strengthened alloys have growth rates at 4K about 4 to 5 times

higher than the lower nitrogen grades. In contrast, the 316 alloy, which

has greater austenite stability than 304, has essentially the same growth

rates at 4K as the 316LN alloy.

3.4 Physical Properties

In general, the physical properties of the N-strengthened stainless

steels are similar to the 300-series alloys as summarized in Table 3.

In general, the nitrogen contributes to austenite stability and thus re- duces the magnetic permeability induced in some alloys by deformation at

low temperatures. Anamolous behavior has been observed in high manganese

466
alloys by Ledbetter (33) in the elastic behavior and by Ledbetter and

Collings (34) in the magnetic susceptibility.

4. Fabrication Characteristics of Austenitic Stainless Steels

The fabricability of austenitic stainless steels is certainly sufficient for fusion magnet systems as attested by the widespread use of

stainless steels in a variety of industries. Fabrication characteristics

that should receive special consideration for fusion magnet applications

are briefly reviewed in this section.

4.1 Welding

The austentitic stainless steels, with and without nitrogen strenthening, are readily weldable by all of the common welding processes, providing the appropriate consummables 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 5. The strength of stainless steel welds at 4K generally exceeds the corresponding base metal

strength, but toughness is usually significantly lower. Three phenomena

that affect the strength and toughness of the as-deposited weld metal

take on added significance at cryogenic temperatures: sensitization,

ferrite content and nitrogen pickup. In addition, heat treatment generally has a stronger influence on the weld metal than the base metal.

Sensitization, the grain boundry precipitation of Cr carbides re- duces weld metal toughness at low temperatures (35). Consequently, weld

metals with extra low carbon, such as 308L and 316L, are commonly used

for cryogenic applications. Sensitization is most extreme if the weldment is given prolonged exposure in the temperature range of 500 to

800C. However, the underlying weld beads in multipass weldments may re- ceive sufficient exposure to these temperatures to cause significant

toughness reductions in stainless steels with greater than 0.03% carbon.

467
Subsequent annealing of the welds at temperatures greater than 950C

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 microfassuring the

weld metal chemistry is generally balanced to provide 4 to 10% ferrite.

However, for cryogenic service, there is concern that the ferrite, which

is a bcc phase and therefore brittle at low temperatures, may embrittle

the weld metal. Szumachowski and Reid (36) have shown that ferrite

reduces toughness at 76K for a wide range of stainless steel weld

metals. Read, McHenry and Steinmeyer (35) have evaluated the effect of

ferrite content on the fracture toughness of shielded metal arc welds of

AWS 316L; the results at 76K and 4K are shown in Figure 9.

Nitrogen increases the yield strength and decreases the toughness of

stainless steel weld metals. When conventional (low N) filler metals are used, nitrogen variations in the range 0.04 to 0.13% may occur depending

on the amount of air which enters the arc shielding medium. The greatest

variations in N-pickup are associated with gas metal arc welding where

the N content ranges from 0.04 to 0.13% depending on the type of gas

coverage and the electrode stickout. Smaller variations are attributed

to electrode coatings; lime coatings generally give better coverage and

less N pickup than titanla coatings.

4.2 Heat Treatment

One of the principal advantages of annealed austenitic stainless

steels is that heat treatment is generally not required. The possible

exceptions are annealing at 1000 to llOOC following severe forming operations and stress relief at 420 to 480C to reduce peak residual stresses

and improve dimensional stability.

Heat treatment of stainless steel welds has been studied by Bennett

and Dillon (37) for AWS types 308, 308L and 310. Heat treating or slow

cooling in the sensitization range causes significant toughness loss at
76K in types 308 and 310, but does not affect 308L. Stress relieving at

840C followed by furance cooling causes about a 50% toughness loss in

each alloy (20-32J as-welded and 8-20J stress-relieved). Stress re- lieving at 950C followed by furnace cooling caused a 50% toughness loss

in 310, no change in 308L and mixed results (+ 40% toughness change) in

308. Annealing at 1070C followed by water quenching improved the toughness of 308 and 308L, but decreased the toughness of 310. The beneficial

effect of a full anneal on toughness at 76K has also been reported by

Krivobok and Thomas (38), McConnell and Brady (39) and DeLong (40).

4.3 Castings

The austenitic stainless steels are widely used for manufacturing

corrosion resistant castings. The casting alloys equivalent to AISl

types 304 and 316 are ACI (Alloy Casting Institute) types CF8 and CF8M

respectively. The ferrite content of these alloys may vary from 0 to 40%

due to chemistry variations within the allowable ranges. Ferrite levels

in this range are likely to be unacceptable because of low fracture

toughness and high permeability. Ferrite-free castings require use of

higher nickel content alloys such as ACI grades CK 20 (corresponding to

AISI grade 310). A 1 6Cr-20Ni-10Mn-2 . 2Mo alloy has been used for liquid

hydrogen (41) and liquid helium bubble chambers (42,43); the alloy re- tains excellent strength (550 MPa yield strength) and ductility at temperatures to 2.5K (43).

4.4 Machining, Forming and Forging

Machining: Annealed austenitic stainless steels are more difficult

to machine than carbon steels or aluminum alloys because of their relatively high tensile strength, work hardening rates, and low thermal con- ductivity. These properties cause poor finishes, heat buildup and excessive tool wear. As a result, it is necessary to use heavier feeds,

slower speeds, better quality tool bits, more rigid machines and more

power.

Forming: The candidate stainless steels have excellent formability.

Alloys with more austenite stability than types 304 or 316 may be needed

if severe forming operations are required. The nitrogen-strengthened

409
21Cr-6Ni-9Mn alloy is reported to have excellent austenite stability; even after 60% reduction at room temperature, the permeability is less

than 1.02 (44),

Forging: The austenitic stainless steels can be forged into shapes

suitable for many magnet applications, e.g. forged pancake billets for

ring rolling coll forms. Stainless steels are more difficult to forge

than most common structural alloys because of their high strength at elevated temperatures. Within the family of austenitic stainless steels,

the forgeability of types 304 and 316 is rated good (45).

4,5 Inspection

During fabrication of fusion magnet systems, the stainless steel

material and components are likely to be inspected using conventional

methods. The only problem peculiar to the proposed designs is the need

for heavy-section stainless steel welds. These welds cannot be

effectively inspected for the presence of cracks or crack-like defects.

Radiography is of limited use because of its inability to reliably detect

cracks. Ultrasonic methods do not work effectively because of the high

noise levels and large and variable attenuation arising from the largegrain-size welds. Standards are needed to assure that the defect sizes

that can be reliably detected do not adversely affect structural

integrity.

5. Aluminum Alloys

Aluminum alloys have been proposed for many structural applications

in fusion magnet systems. Their princhpal advantages are: low as-fabricated cost, light weight, non-magnetic behavior, stable microstructure,

and good retention of strength and toughness at cryogenic temperatures.

The main disadvantages of aluminum alloys are lov/ strength in weldments,

high electrical and thermal conductivity and an unfavorable elastic modulus (too low) and thermal expansion (too high) for many applications.

470
For fusion magnet structures, the preferred aluminum alloys are

5083, 2219 and 6061, The composition ranges for these alloys are summarized in Table 6. Alloy 5083 is generally used in the annealed condition

(5083-0) and has a typical yield strength of 171 MPa (25 ksi) at 4K.

Alloy 2219 is generally used in a heat treated condition (T 851, T 87

etc.), and in the T 851 temper has a typical yield strength of 482 MPa

(70 ksi) at 4K. Alloy 6061 is usually used in the heat treated condition

(T 6, T 651, etc.), and in the T-6 temper has a typical yield strength of

380 MPa (55ksi) at 4K. Although each of these alloys is readily weldable, only 5083 retains full strength in the welded condition and is the

best choice for most large welded components. Alloy 2219 has the highest

strength of the three alloys and is recommended for bolted construction

where space or weight need to be conserved.

Several other aluminum alloys have excellent properties at 4K and

may be used for special applications. Commercially pure aluminum (alloy

1100) and the 3003 alloy have excellent formability. Alloy 3003 has been

widely used for cryogenic applications and is the best choice when brazing is used for joining. The high strength alloys such as 2014 and 7005

may be used for nuts and bolts.

The ASTM specifications for aluminum mill products of alloys 5083,

2219 and 6061 are summarized in Table 7. For a more complete listing of

specifications refer to the Aluminum Standards and Data Book (46). Procurement of heavy sections may require additional specification requirements on through-thickness ductility (47).

The low temperature properties of aluminum alloys have been compiled

in cryogenic handbooks (6,7), in a review article by Kaufman and Wanderer

(48) and in suppliers publications (49). In this section, the general

trends in strength, toughness and fatigue resistance are summarized,

particularly the data for room temperature and 4K. The data are taken on relatively few lots of material, but generally the lot-to-lot variations

in the mechanical properties of aluminum alloys are relatively small

(with respect to stainless steel).

471
5.1 Tensile Properties

The mechanical properties of alloys 5083, 6061 and 2219 at room temperature and 4K have been determined by Kaufman, Nelson and Wanderer

(50) and the results are summarized in Table 8, Notice that in all cases

the tensile strengths, and to a lesser extent the yield strengths are superior at 4K and the ductility values are not significantly changed.

The alloys retain good notched-tensile properties at 4K.

Aluminum alloys having a wide range of strength levels are suitable

for liquid helium applications. The yield strengths of commercially pure

(CP) aluminum (types 1100-0 and 1099-H14) and the preferred alloys are shown as a function of temperature in Figure 10 (42). The yield

strengths of each alloy gradually increase with decreasing temperature.

Strength increases in the following order: annealed CP (1100-0), cold

worked CP (1099-H14), solution-strengthened Al-Mg alloy (5083-0), cold

worked Al-Mg alloy (5083-H321), precipitation hardened Al-Mg-Si alloy

(6061-T651) and precipitation hardened Al-Cu alloy (2219-T851). Still

higher yield strengths are available in several other alloys including

2014-T651, 2024-T851 and 7005-T5351.

5.2 Toughness

Aluminum alloys generally retain excellent toughness at cryogenic

temperatures. A useful measure of aluminum alloys toughness at 4K is the

notch-yield ratio, i.e., the ratio of notched tensile strength to un- notched yield strength (51). Ratios substantially greater than one, say

1.5, indicate that the alloys retain the ability to plastically deform

and resist crack initiation in the presence of sharp stress concentrations. As shown in Table 9, the 5083 and 6061 alloys have notch-yield

ratios greater than 1.6 and the higher-strength 2219-T851 and T87 alloys

have notch-yield ratios of about 1.4. Significantly lower ratios (1.0 to

1.25) occur in other high strength alloys such as 2014-T651 and 2024-T851

and in most 7000 series alloys (7005 is an exception), where the ratio is

sometimes below 0.4 at cryogenic temperatures (52).
Limited toughness test data are available on aluminum alloys at 4K.

Tobler and Reed (53) measured the fracture toughness of 5083-0 (4.3 cm thick plate) using J-integral procedures and found that toughness increased with decreasing temperature. The Kj^(J) was 47 MPa/iri at 4K

in the transverse (TL) orientation. Shepic (54) measured the toughness

of 2219-T87 (3.8cm thick plate) using ASTM E399 plane strain fracture

toughness testing procedures. The for 2219-T87 at 4K was 45.5

MPa/m in the transverse (TL), 47.7 MPa/in in the longitudinal (LT) and

28.6 MPa/m in the through-thickness (SL) orientations.

5.3 Fatigue

Strain cycling fatigue tests have been conducted by Nachtigal (13) on 2219-T851 and 2014-T6 alloys at room temperature, 77K and 4K. For a given strain level, the fatigue lives of both alloys were greater at 4K

than at the higher temperatures for endurances greater than about 500

cycles. The results for 2219-T851 are shown in Figure 11. The MansonHirschberg (55) method of universal slopes gave reasonable predictions of

fatigue behavior based on the tensile properties at the test temperature.

The fatigue crack growth behavior of 5083-0 (53), 2219-T87 (54) and

an Al-6Mg (56) alloy have been determined at 4K. In each case, the re- sults indicate that the growth rates at 4K are slower than the rates at room temperature.

5.4 Physical Properties

The physical properties of the candiate aluminum alloys are summarized in Table 9. The thermal properties have been taken from the bestfit lines used in the LNG Materials Handbook (7) and from the compilation

on thermal conductivity by Childs, Ericks and Powell (57). Original

sources are given for the elastic properties (58,59) and electrical re- sistivity (17).

With tlie exception of magnetic susceptibility, the physical properties of aluminum alloys are not as good as those of stainless steels for

473
fusion magnet applications at 4K. The elastic modulus is relatively low,

and thus when used for coil case applications, more of the load must be

carried by the conductor assembly. The thermal and electrical conductivities are relatively high, resulting in incrased thermal loads particularly from eddy current losses.

6. Fabrication Characteristics of Aluminum Alloys

The relative ease of fabricating large aluminum structures is the

principal advantage of aluminum over stainless steels. In addition the

lower weight of aluminum structures may facilitate assembly and erection

of the large magnet structures anticipated for fusion energy systems.

Fabrication characteristics that should receive special consideration for

fusion magnet systems are briefly reviewed in this section.

6.1 Welding

Alloys 5083, 6061 and to a lesser extent 2219 are readily weldable

by the gas metal-arc and gas tungsten-arc processes. The compositions of

aluminum welding rods and electrodes for these processes are summarized

in Table 10. Annealed 5083 alloy retains full strength in the as-welded

condition, but the cold worked and precipitation hardened alloys are weakened by welding. Part of the strength of precipitation hardened

alloys can be restored by a post weld heat treatment.

The tensile properties of aluminum alloy weldments were determined

by Nelson, Kaufman and Wanderer (60) and the results for selected alloys

at room temperature and 4K are summarized in Table 11. Notice that part

of the strength obtained by cold working (5083-H321) and by precipitation

hardening (6061-T6 and 2219-T851) is retained by the welded joint. This

strength increase is not reflected in design-code allowable stresses (11)

which assume as-welded aluminum alloys are in the fully annealed condition. The tensile properties of 6061-T6 and 2219-T62 welds that were subsequently solution treated and aged indicate that substantial

strengthening can be achieved by post-weld heat treatments. For most

fusion magnet structures, the solution treatment and quenching are not

considered practical. The data in Table 11 also indicate that the

474
ductility of 2219 weldments is quite low due to localized deformation in

the heat affected zone; the ductility and notch-yield ratio of 5556 and

4043 filler metals (as-welded) are low at 4K; and the best ductility and

notch-yield ratios in the as-welded condition at 4K are provided by the

5083 and 5356 alloys.

6.2 Heat Treatment and Temper Designations

Aluminum alloys are supplied in a wide range of tempers, and for

large complex structures, subsequent heat treatments are not generally

performed by the fabricator. The tempers are designated by a letter

followed by one to three digits. The letter is the basic temper designation: 0-annealed, H-strain hardened and T-thermally treated. For the

strain hardened tempers, the first digit indicates the hardening operations: Hl-strain hardening only, H2-strain hardened and partially an- nealed and H3-strain hardened and stabilized. The second digit (e.g. H18

is strain-hardened by a cold reduction of 75%) indicates the degree of

strain hardening and the third digit is sometimes used for specified

variations. For the thermally treated tempers, the first digit indicates

the specific sequence of basic treatments; for example T6 designates

solution heat treated and then artificially aged and T8 designates solution heat treated, cold worked and then artificially aged. Subsequent

digits indicate specific variations, such as x51 stress relieved by

stretching as in T651 or T851. The reader is referred to the Aluminum

Association Handbook (46) for a complete description of the temper designation system.

6.3 Castings

Aluminum alloy castings can be produced by a variety of processes

and generally are satisfactory for non-structural applications at low

temperatures. The ASME Code (11) permits use of Al-5 Si (Type 43) and

Al-7Si-0.3Mg (Type 356-T6) for non-welded structural applications.

Nelson, Kaufman and Wanderer (60) evaluated the tensile properties at 4K

of several combinations of castings welded to wrought alloys; in general,

superior properties were obtained when the properties of the cast and

wrought components were nearly equal.

475
6.4 Machining, Forming and Forging

Machining: Aluminum alloys, particuarly the heat treated grades,

have excellent machinabillty . Machined and bolted assemblies of high

strength aluminum alloys can replace more costly welded stainless steel

assemblies. Care should be taken when using steel bolts because the

thermal contraction of aluminum alloys on cooling to 4K exceeds that of

steel.

Forming: Aluminum alloys have excellent formability in the annealed

or the solution treated conditions.

Forging: Aluminum alloys have the best forgeability of all structural alloys because of their very low strength at forging temperatures

of 370 to 510°C. The size limit on aluminum forgings is significantly

greater than that on stainless steels because of the reduced pressures

required

.

6.5 Inspection

Aluminum alloys and their weldments are readily Inspectable by con- ventional methods: dye-penetrant, ultrasonics and radiography.

7. Alternative Alloys

Alternative alloy systems do not provide candidate structural materials for superconducting magnets in fusion energy systems. Consider the

following systems:

Nickel base alloys are ferromagnetic, expensive and frequently difficult to fabricate. However, many of these alloys have excellent com- binations. of strength and toughness at 4K and may be suitable for special applications (61).

Titanium alloys are brittle at 4K, expensive and difficult to fabricate (62).

Cobalt base superalloys transmute to long half-life reaction products and are expensive and difficult to fabricate.

47G
Carbon and alloy steels including the ferritic and martensitic stainless -

steels are ferromagnetic and brittle at 4K. Even the cryogenic nickel steels

are brittle at 4K (63) except for a specially processed 12Ni alloy (64).

Copper alloys are acceptable, but generally not competitive with the aus- tenitic stainless steels or aluminum alloys.

8. Summary Comments

This review suggests that there are several commercially available alloysthat are suitable for construction of large superconducting magnet structures

for fusion energy systems. These alloys include annealed austenitic stainless

steels, nitrogen strengthened austenitic stainless steels where greater

strength is needed, an annealed Al-Mg alloy for welded aluminum construction

and precipitation-hardened aluminum alloys for bolted assemblies. For each of

these alloys, there are enough experimental data and fabrication experience to

provide reasonable assurance that the alloys have sufficient strength, ductility and toughness in the liquid helium environment; that the alloys can be fabricated into support structures for superconducting magnets; and finally, that

the fabricated structure will retain adequate properties during operation of

the fusion energy device.

477
It should be realized, however, that the depth of understanding, the

degree of characterization, and the extent of related experience for these

alloys in the fusion-magnet environment are extremely limited. More work is

needed on material behavior at 4K, particularly in the areas of material variability, processing effects on properties, and the fabrication and mechanical

behavior of large, complex, thick-section components.

9. Recommendations

The accumulated knowledge of materials for fusion energy systems should

eventually result in a code for the design, construction and inspection of

fusion magnets. Adherence to this code should provide assurance that the magnet system provides the intended field and operates in a safe, reliable and

efficient manner. Establishment of such a code for large magnets must await

further developments in the technology, particularly the experience gained from

prototype systems and greater material characterization. However, parts of the

code should be set forth as preliminary guidelines to facilitate the orderly

development of materials technology for magnet systems.

Design guidelines needed for selection of materials should Identify operating loads and conditions and criteria for establishing allowable working

stresses for candidate materials. For magnet systems, description of the operating loads should include the anticipated load history for normal operation,

loads due to credible fault conditions and procedures for treating combined

loads. The operating conditions of interest are the thermal environment, radiation dosage, magnetic field profile and design life. Allowable working

stresses should account for the materials strength variability, the fatigue

strength for a specified scatter factor on fatigue cycles, fracture mechanics

analysis and a factor of safety based on uncertainties in the loads and stress

analysis

.

Material procurement and processing should be controlled by a set of specifications which are sufficient to preclude use of materials in the magnet

systems which have properties inferior to those assumed in design. Inspection

specifications and standards are needed to provide a high degree of confidence

that no defects are present in the raw materials, or introduced during fabrication, which could cause unsatisfactory performance of the magnet system.

478
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479
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Stahl and Eisen, Vol. 91 (1971) p. 1255.

22. R. Voyer and L. Weil, in: Advances in Cryogenic Engineering, Vol. 11,

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23. D. C. Larbalestier , Private Communication, U. of Wisconsin, Madison,

WI (1977).

24. C. E. Spaeder, Jr., and W. F. Domis , Cryogenic and ElevatedTemperature Mechanical Properties of a High-Nitrogen Type 304 Stainless

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

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4U0
29.

P. F. Scardigno, M. S. Thesis, Naval Postgraduate School, Monterey,

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255.

34.

H. M. Ledbetter and E. W. Collings, in: Metal Physics of Stainless

Steels, TMS AIME, New York, Forthcoming.

35.

D. T. Read, H. I. McHenry and P. A. Steinraeyer, in: Materials

Studies for Magnetic Fusion Energy Low Temperature Applications - II,

National Bureau of Standards, Boulder, CO (1979).

36. E. R. Szumackowski and H. F. Reid, Welding Journal, Vol

325-s.

37. V. N. Krivobok and R. D. Thomas, Jr., Welding Journal,

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39. J. H. McConnell and R. R. Brady, Chemical Engineering,

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. 57 (1978) p.

Vo. 29, (1950)

Vo. 67 (1960)

Vol. 20,

44.

Armco Nitronic 40 Stainless Steel Sheet and Strip, Armco Steel Corp.,

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Solids at Room Temperature and Below, NBS Monograph 131, National Bureau

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Cryogenic Engineering, Vol. 14, Plenum Press, New York (1969) p. 71.

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Society for Testing and Materials, Philadelphia, PA (1976) p. 346.

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(1975) p. 1569.
Table 1 - Compositions of Austenitic Stainless Steels

Suitable for Liquid Helium Service

AISl

Type No. Cr

304 18-20

304L 18-20

316 16-18

316L 16-18

310 Ik-lh

310S 24-26

For each grade Mn = 2.0 max

Si = 1.0 max

Composition, %

Ni C ,max Other

8-10.5 .08

8-12 .03

10-14 .08 2-3 Mo

19-22 .03 2-3 Mo

19-22 .25 1.5 Si

19-22 .08 1.5 Si

= .030 max

= .045 max

max

max

483
Table 2, ASTM Specifications for Various Product Forms of Austenitic

Stainless Steels

Product Form ASTM Specifications

Plates, sheet and strip A240

Forgings A473

Bars and sections A479

Seamless tubes A213

Welded tubes A249

Seamless and welded tubes A269

Small diameter tubing A632

Bolts, screws, studs A320

Nuts A194

484
Table

3,

Physical

Properties

of

Austenitic

Stainless

Steels

NO NO r>- 00

OJ 1—1

rH 1—1

f—

H

rH rH

oC 00 CN

CN a^

O • r—( • • f—t o CO cn iH rn >c CO •-H g m CN ON S-( M <u NO NO

C/i JC .

t M H CO CO < 00 CO

ON NO T3 03 o o CN 1—1 C C o o m • • CO CO CO • • 00 o m iH CM CO OD O O r*-. CN 00 C .—1

f—

H

03 CO CO O

hT

CO

O o

CO CO

M-l CTn vO vO VnD MT 5-1 00

(U T—

H

rH 1—1

•—( O

O o I—1

rH

PC CO C4H C4H

CD QJ QJ o e 6 E

CT3 CO CO

m CTn CO CO CO o CN m rH M • OJ QJ 0) • • CO Mj- C7> O Z z 43 00

h-1 CTn 4-1 4H 4H VsO < •—t

>, So cn o rH rH 1—H o On cn >-4 Sj ^4 CO o • • CO CO CO • • UO O cu d) QJ rH o^ 00 O z z Z 00

CM

4-1 ON vD NO v£> O' 00

(U —H

PC

O

COM C3N CO CN CO CNJ m NO O hH o • < 1 o o o o o CT^ iH 1 CN 00 rH oON CN

CN oo o

S4: • • • • • m C3N o CO m o o CJN 00 r-H i-H 00

CN •—1 Io M

4-) •Hc

1=3

CO CN CM Bee

00 o o e • • CJ 00 Z cj I-) eu

CJ3

CO o 3 •H

f-H CO 4-1 3 3 CO X) ,

1 oi o 3 z T3 cn So Q —

4H So CO c

S-I 4-4

— o

QJ •H 00 cn O. CO c CQ cn O c 3 QJ •H

>-i QJ o X

O

CL,

Q s>-< CO CU

.u >>

•H 4^

4-1 > •H

•H •H fH > Ui •H

•fH cn

X 4H •H CO

CJ cn QJ 3 4H <v e

03 CO 0^u

C QJ QJ ^ O Z iH CU rH

CJ

C CO CO O U CJ O -H rH 'H •H •H •H iJ

CO CO >4H

U 4J -H e c •H 4-1 QJ C u CO O CJ C HH

QJ CL QJ 0) 00 Nr X X CL _C0 H u CO CO Z

‘rob
Table 4. Compositions of Nitrogen-Strengthened Austenitic

Stainless Steels Suitable for Liquid Helium Service

Type Nominal Composition, %

Cr Ni Mn C, max N Other

AISl 304N 18-20 8-10.5 2.0, max .08 .10-. 16 - 304LN 18-20 8-12 2.0, max .03 .10-. 16 - AISI 316N 16-18 10-14 2.0, max .08 .10-. 16 - 316LN 16-18 10-14 2.0, max .03 .10-. 16 - ASTM XMll* 18-21 5-7 8-10 .04 .15-. 40 - ASTM XM19* 20.5-23.5 11.5-13.5 4-6 .06 .20-. 40 Cb .10-.

V .10-. 30

Mo 1. 5-3.0

*ASTM A412
Table 5 - Compositions of Filler Metals for Welding

Austenitic Stainless Steels

AWS Type Composition, %

Covered Electrodes Cr Ni C, max Other

E308 18-21 9-11 .08

E308L 18-21 9-11 .04

E310 25-28 20-22.5 .20 .75 Si, max E16-8-2 14.5-16.5 7. 5-9.5 .10 .50 Si, max; 1-2 Mo

E316 17-20 11-14 .08 2-2. 5Mo

E316L 17-20 11-14 .04 2-2. 5Mo

E330 14-17 33-37 .25

Rod and Bare Electrodes

ER308 19.5-22 9-11 .08

ER308L 19.5-22 9-11 .03

ER310 25-28 20-22.5 .08-. 15

ER316 18-20 11-14 .08 2-3MO

ER316L 18-20 11-14 .03 2-3Mo

For Covered Electrodes: Mn = 2.5, max; Si = 0.9, max; P = .04, max; S = .03 max

For Rod and Bare Electrodes: Mn = 1.0-2. 5; Si = .25-60; P = .03, max; S = .03, max

4G7
Table

6.

Compositions

of

Aluminum

Alloys

Suitable

for

Liquid

Helium

Service

(46) eDBcB

CO • Pi rH > m

CU < CN o • o 1 I O • m o

c:^ o 1^

la^ • • X

X o ca CO -H EE • ^ I uo m

•H I CM I 1—( —( H I o I • • X C ct3 N1 B O O• • o o o LO cn

.H CM CM o o o u I I m m

CM

m• • I I U~1 ^ O O

6^

•> t>o c s o

•1—I X

COB

CMO

CJ^ CM • • I—I I I O C30 I • o t MT .

Cfi oaBo c

cj s X o CO <r B • " I LO O O CM

cn o o o• 'O' 3a O CO

CM • • ^I I in CO o • • u~i

O

X o CM CO Mf • B • I » I m iH cn O • . o •

QJ

X

pLi cOB

X

CO CO B • ^ o

OJ +

p'' <r • • • o o o

•H

X

CO COB

O CM • • --I o

C30 \D 'O' • • • I O O 'O'

QJ o on CO

O' o O 00

•—1

<>4 o O

H r-H CN CO uO vO 4C8
Table 7, ASTM Specifications and Availability of Aluminum Mill Products (46)

Product Form Specification Availble Alloys

ASTM 5083 6061 2219

Sheet and plate B209 X X X

Wire, rod, bar; rolled or B211 X X

cold finished

Wire, rod, bar , shapes , tube B221 X X X

Structural shapes B308 X

Pipe B241 X

Forgings and forging stock B24/ X X X

Structural pipe and tube; B249 X

extruded

4G9
Table

8,

Tensile

Properties

of

Aluminuin

Alloys

at

Room

Temperature

and

4K

(50)

IT) CM o CM CO oo NO oo oo o 00 r-- o o cn cn nD <r 00 00 NO NO v£) NO UO cn cn un H >-< Z CM CM 1—1 —1 rH r—1 rH rH t—

I 1 1

—

rH rH 1—

NO NO NO 00 CM NO NO 1— CN cn CO

C/0 CM 00 to CM LT) CN 1—

o rH o CN o CM o

H C/i z H iH o 1^ o 1—

f-H 1—i 1—

1—

iH rH —1 iH

01 x:

CO Zu

iH

•H

4-1 OO < CM •-H 00 CJN r-- CO un o Nt un H 4-1 (/) c Oh CN CM o r-- r-H o <r O 00 NO NO CJN m a> Z o c <u S CO <r in NO NO UO UO NO LT) NO cn z 0) H

(-1 4-1 cn

U Cfl

•X} '+-1 O) o

Pi

O *H

6^ X r~>

CNJ CO

CO CO

CNJ CO O CNI CN! COuo -<r Nj- CO

r-~ NO CNI ^

CN CM CNI CN 00 COCN CN O CTN

4-1 Q UO UO un CM CM 00 00 iH

CC -<r • • • • • • • • 00 S'? ON CN UO o^ MD UO UO CO un O CO »H m 00 c C CO r—

CN rH CM «H CN iH rH •—

o •H

T—

w

X4-J xl OO

1—1 c cO f-H UO i—j o> 00 CO CO r-* i-H 00

CL) 0) Oh CO On v£> CO UO 00 00 o 00

•H 1-4 Z •—

p—

(

CM CM CN CO CN CO m CO <r CO UO CN CO

4-J cn

<U z

I—1 4-1 H 00 03 CM UO 1

CO CT\ UO NO C7^ r-- CJ> UO

cn c Ph CM UO CO O 00 O CXD NO UO UO UO r-- CO UO C CD S CO UO CO UO CO <r CO <r --a" o vD o CN

OJ S-i H 4J cn

D-B z H H H H H H H H

CD 0^ Pi P< PP Pi

CO

CO

0) OJ •U C

CD Z e UO UO UO UO un o^

i-l O

Ph -H B CN CM CM CM CN fH X

H

T3 ?—

i“H C CM LO UO >—

CD Lh CO 'X) 00 00 MO tiO

0) O PP H H H H c

>N O- 1 ! 1 1 •H o e CO CO 1—

a^ CON MO 4-J 1—1 (U 00 oo X> f—

(

r~( UO CO o O CM CN CO cc3 u‘i VsD C^4 CM < O

490
Precipitation

Hardened

Precipitation

Hardened

5083

(Annealed)

6061

2219

Property

Units

295K

4K

Ref

295K

4K

Ref

295K

4K

Ref ON o^ a^ o m m m rH CO •

H v£) oo cn 1 CM 05

• • • • • • 85 32 ro CN

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r—

H

1—

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m i-H

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CN CO 1 —1 CN m

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00

CO o o

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120 120 120 CO

00 00 00

u-i LO

00

00 CO

ON CO CO —

<

CN o

o o CO cn

00 ro •-H

CO

50 CO 50

50 LO 00 CO 50

• • • • • CN o cn o tn CN CN CN o •—

1

<Tv

NO

1o

cn CN CN »—

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0) 3 (U H <u CU w (0 0^ QJ D. o JO o J= JO O- »—

<

CO >1 CO Cl, H H CO w

491
Table

10.

Compositions

of

Aluminum

Welding

Rods

and

Electrodes

for

Liquid

Helium

Service

(46)

bO

'V

C CO I-H f-H f-H 01 •H o NO NO

A-i o o O

O

CO CO NO NO

OJ 0) • #1 &o s o <J\ o CO CO CO o f-H o 00 oo oo d Xi »—1

CN r—{ o o o

cn o t—jCN UO u-i LO

f—1

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x: • 1—1 o 4-» co^ • • o <d> o o

00 00 00 00 oo o o o o o

OJ o o o o o

PQ o o o o o o o o o o o o o CN CN CN • o UO • • o CN 1—H o o

•H 1 • • 1 1 H o o o vd UO

1—( o o • • • o o o o o m o UO

C •—

H

I-H CN ^H CN

N • • • • • o o o o o o o o X CO 00 X X

CU CO • • • 03 cC

U-i E o o o E E c

0) o UO > • • • •H 1—1 o o o

•H II o • o II II C/!) 0) CN <r 01 o;

<U • 1 • pH tU| M + o UO o + + C •H • •H •H

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Vj

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<u • • •

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6^ o o

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4-J CN cn o • •H • • • o

CO o o o o

cxE CN UO UO o CN cn • • • u 00 1 o o UO UO LO s 1 • • 1 1 1 o o CO lO • • •

00 o • CO o o o d CNJ o f-H ^H 1—H

Cd • 1 • • • • o 00 o o o o

UO

0) o <T> CO CO NO vd

(X o 00 UO CO o I-H CO LO H CN m UO m

492

V*
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rH r—

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tn 00 CO C7N o o CM CM 1 00 <r 00 rH LO

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CO 00o H 4J CO CU Z CO CO CO Z o c p z (U X H c/a ol£) (UP z C 3

MT

O

X

*iH •H IN XI 4-1 CO PQ <£ c CO iL PC

CO CJO

X

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!~i 3

4-1 >L

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<u 4-1 (U 0 CT^ MO a d •H 0 vO ON e •tH U 6M r-H (V O •rH H X X e w o d o o

pi •H CO 4H (U 4J CJ p 1 Os

CO 3 c SM 1 CO CO X X CU X

4-1 pi o c

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4-1

Q LO

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pi

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CN

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r*-.

m

CN cn

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0)

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CN

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X

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493

Note:

Location

of

Failures

W1-

Weld

metal

BMBase

metal

HAZHeat

affected

zone
List of Figures

1. Tensile and yield strengths of three austenitic stainless steels - AISI types 304, 310 and 316 - at temperatures between 4K and 300K

( 6 ).

2. Tensile properties of two austenitic stainless steels - AISI types

304 and 304L - at temperatures between 4K and 300K(6).

3. Notched tensile properties of five austenitic stainless steels - AISI types 304, 304L, 310,310S and 316 - at 4K (6). K'p is the

stress concentration factor of the notch.

4. Fatigue behavior of three austenitic stainless steels - AISI types

304 and 316 and a 21Cr-6Ni-9Mn alloy - at 4K (14),

5. Fatigue crack growth rates of AISI type 304L austenitic stainless

steel at 295, 76 and 4K (15).

6. Tensile and yield strengths of nitrogen-strengthened 300-series

stainless steels at temperatures between 4K and 300K (20-27).

7. Tensile and yield strengths of 21Cr-6Ni-9Mn stainless steel at

temperatures between 4K and 300K (28).

8. Fracture toughness of four nitrogen-strengthened austentitic

stainless steels at temperatures between 4K and 300K (25).

9. Fracture toughness of AWS 316L weld metal at 4K as a function of

ferrite content (35).

10. Yield strengths of six aluminum alloys at temperatures between 4K

and 300K (50).

11. Fatigue behavior of 2219-T851 aluminum alloy at temperatures of 300,

77 and 4K (13),

494
STRENGTH,

MPa

1. Tensile and yield strengths of three

austenitic stainless steels - AISI types

304, 310, and 316 - at temperatures between

4K and 300K (6).

496
STRENGTH,

MPa

1800

1600

1400 -

1200 -

1000 -

800 -

600 -

SHEET

I Ultimate Tensi le

•304LJ Strength

100 200

TEMPERATURE, K

2. Tensile properties of two austenitic

stainless steels - AISI types 304 and 304L - at temperatures between 4K and 300K (6).

496
CJ3

CdI—

ooQ

>-QLUO

h-O

Ozr

LU

a:I—

oo

00

QLU

C_) I—O

n=I—

C2

LU

q;HOO

oo

(_! I—o

o

LU

cc: I—

oo

oo

Q

LU

rc

C_) I—o

3. Notched tensile properties of five

austenitic stainless steels - AISI types

304, 304L, 310, 310S, and 316 - at

4K (6). Kj is the stress concentration

factor of the notch.

497
TOTAL

LONGITUDINAL

STRAIN

AMPLITUDE,

percent

4. Fatigue behavior of three austenitic

stainless steels - AISI tyoes 304 and 316 and

a 21Cr-6Mi-9Mn alloy - at 4K (14).

r

493
FATIGUE

CRACK

GROWTH

RATE,

da/dN

(mm/

cycle)

10 100

STRESS INTENSITY FACTOR RANGE, AK, MPa-m"

5. Fatigue crack growth rates of AISI type

304L austenitic stainless steel at 295, 76

and 4K (15).

400
STRENGTH,

MPa

6. Tensile and yield strengths of nitrogenstrengthened 300-series stainless steels at

temperatures between 4K and 300K (20-27).

GOO
STRESS,

GPa

7. Tensile and yield strengths of

21 Cr-6fli-9Mn stainless steel at temperatures between 4K and 300K (28).

SOI
3. Fracture toughness of four nitrogenstrengthened austenitic stainless steels at

temperatures betv/een 4K and 300K (25).
FRACTURE

TOUGHNESS,

(MPa/m)

9. Fracture toughness of AWS 316L weld

metal at 4K as a function of ferrite

content (35).

503
YIELD

STRENGTH,

MPa

10. Yield strengths of six aluminum alloys

at temperatures between 4K and 300K (50).

504
TOTAL

LONGITUDINAL

STRAIN

11. Fatigue behavior of 2219-T851 aluminum

alloy at temperatures of 300, 77 and 4K (13).

SOS

HANDBOOK PREPARATION

Mechanical Properties Data Center

National Bureau of Standards

j07

HANDBOOK PREPARATION

F. R. Fickett

Thermophysical Properties Division

National Bureau of Standards

The handbook is being produced jointly by the Mechanical Properties

Data Center (MPDC) of Traverse City, MI and the staff of the NBS-Thermophysical Properties Division. The literature searching and document acquisition is done by NBS. MPDC extracts the data and prepares worksheet graphs

and characterization tables which are then evaluated by the NBS staff to

determine "best value" curves and appropriate error bands. The handbook

presents "best value" curves for properties in two catagories listed in

Table 1. For the category A properties we give an error band which indicates the spread that one might expect due to material variability, experimental error, etc. A typical graph is shown in Fig. 1. The published graphs

will have SI units primary, but common units will be shown also. Each A

category property presentation also contains a characterization table giving

the appropriate information (different for physical and mechanical properties)

on each of the original data sources used in constructing the best value

curve. A preliminary version of the table is shown as Table 2. Another

table is given for the "A" properties which provides a tabular listing of

the best value curve (see Table 3). The B category properties are presented

in smaller graphs as shown in Fig. 2. Error and variability estimates are also shown for these properties where possible. For all properties a com- plete bibl iograpiy i s included. Each chapter (one for every material) starts

with a detailed discussion of the material including: composition limits,

designations, product forms, heat treatment, melting and casting, fabrication,

welding, and chemical properties.

The materials being processed for inclusion in the first edition of the

handbook are: stainless steels - 316, 304, 21/6/9 (and L and N modifications)

and aluminum alloys - 5083, 2219 and 6061.

We have completed and evaluated the final draft of the first chapter

(AISI 316) and it will be printed soon. The handbook covers have been

received and our plan at present is to start distribution as soon as the

first chapter is ready. Computer runs and evaluation of the output have been

done for the second chapter (21/6/9). The data for the other chapters has

been extracted from the easily-obtainable literature and entered into the

computer system at MPDC. We are now evaluating more obscure documents and

reports prior to doing the initial computer runs on these materials.
Table 1. Properties presented in the handbook. A given property is

included only if low temperature data exist.

A (primary) B (secondary)

Coefficient of thermal expansion Bulk modulus

Crack growth rate Compression

Fatigue strength Creep

Fracture toughness Electrical resistivity

Tensile strength Impact strength

Tensile yield strength Magnetic permeability

Thermal conductivity Poisson's ratio

Young's modulus Shear modulus

Specific heat capacity

Stress-strain

Tensile elongation

Tensile reduction in a

Thermal diffusivity
STRESS

-

MPa

MATERIALS FOR

SUPERCONDUCTING

MAGNET SYSTEMS

AlSl 316

Structural Alloys

Stainless Steels

TEMP - F

-460 -400 -300 -200 -100 0 80

THE YIELD STRENGTH OF ANNEALED TYPE 316, 316L, AND

316LN STAINLESS STEEL (1. 2, 5. 6, 7. 8. 9. 10, 11, 12)

MECHANICAL PROPERTIES DATA CENTER

Figure 1. A typical handbook graph - A-category property. Colored 2.03

shading is used in the actual handbook.
Structural Alloys

Stainless Steels AlSl 316

MATERIALS FOR

SUPERCONDUCTING

MAGNET SYSTEMS

REFERENCE NO. G 7 8 9 .. 11 12

Composition 316 316 316 316LN 316LN 316LN 316L CR 17.52 17.47 17.70 17.73 17.98 17. 8 NT 13.48 13.40 13.63 11.70 13.08 13.9 MO 2. 34 2.28 2.78 2.55 2.77 2.3 MN 1. 86 1.67 0.86 1.35 1.63 1.96

SI 0.58 0.45 0. 62 0.33 0. 20 0.41 N 0.03 - 0. 15 0.163 0.17 - C 0.057 0.058 0.023 0.018 0.027 0.027 P 0.024 0.018 0.018 - - - s 0.019 0.022 0.020 - - - FE Bal. Bal. Bal. Bal. Bal. Bal. Others . 02 TI, .02 CO .11 CU

Heat Treatment .Annealed Annealed Annealed Annealed Annealed Annealed Annealed

Average Grain ASTM 5 Size

Hardness R 79 B

Product Form Round Bar Sheet

0. 1016 cm

thick

Rod Forging

Specimen T^TDe Round Flat Round Button Round (b) Min. Width or 0. 381 cm dia. 0. 50S cm Head 0. 564 cm dia. Dia. Wide

Thickness Reduced Sec.Lg.

0. 1016 cm

3. 5 cm

Gage Length 2. 54 cm 5. 08 cm 5D

Notch Type

Included Angle

Root Radius KT Factor

Strain Rate 0. 0508 cm/min. 0. 01 cm ' cm/ 0. 05 cm/min. min. &

254 cm/min.

Total Spec. No. 3 6 7 “ 1

7 2 5 2

Remarks

(b) Specimens taken from various locations of the forging and with different orientations.

CHARACTERIZATION TABLE FOR THE TENSILE PROPERTIES

OF ANNEALED TYPE 316 STAINLESS STEEL

MECHANICAL PROPERTIES DATA CErTTER_in^

2. 00.1 Table 2. A small part of the character! zati on table for the tensile

properties of AISI 316.
Structural Alloys

Stainless Steels AlSl 316

MATERIALS FOR

SUPERCONDUCTING

MAGNET SYSTEMS

TABULAR PRESENTATION OF THE “BEST VALUE CURVE” FOR THE YIELD STRENGTH

OF ANNEALED TYPE 316 STAINLESS STEEL

NUMBER OF DATA POINTS: 35 TOLERANCE FACTOR A: 259.5 TOLERANCE FACTOR B: 149.8

Temp Temp Temp Temp

K Of MPa KSI °F K KSI MPa

4.2 (-452.1) 657. 8 (95. 41) -452 ( 4.2) 96. 84 (657.7)

20 (-423.7) 636.9 (92. 37) -420 ( 22.04) 91.98 (634.2)

40 (-387.7) 610. 3 (88. 52) -390 ( 38.71) 88.76 (612.0)

GO (-351.7) 583. 8 (84.67) -360 ( 55.37) 85. 56 (589.9)

77 (-321.7) 5G1. 2 (81.39) -330 ( 72.04) 82.35 (567. 8)

80 (-315.7) 537. 2 (80. 81) -320 ( 77.59) 81.28 (560.4)

100 (-279.7) 530.7 (76.97) -300 ( 88.71) 79. 15 (545. 7)

120 (-243.7) 504.2 (73.13) -270 (105.4) 75.94 (523.6)

140 (-207.7) 477. 6 (69.27) -240 (122.1) 72.74 (501.5)

ICO (-171.7) 451.1 (65. 43) -210 (138.7) 69.52 (479.3)

180 (-135.7) 424. 5 (61. 57) -180 (155.4) 66. 31 (457.2)

200 (- 99. G7) 398.0 (57.72) -150 (172.0) 63.11 (435.1)

220 (- 63. G7) 371.5 (53. 88) -120 (188.7) 59.90 (413.0)

240 (- 27.67) 344.9 (50.02) - 90 (205.4) 56. 69 (390.9)

2G0 ( 8.33) 318.4 (46.18) - 60 (222.1) 53.49 (368.8)

273 ( 31.7) 301.1 (43.67) - 30 (238.7) 50. 27 (346.6)

2S0 ( 44.33) 291.8 (42.32) 0 (255.4) 47.06 (324.5)

300 ( 80.33) 265.3 (38.48) 30 (272.0) 43.86 (302.4)

GO (288.7) 40. 65 (280.3)

80 (299.8) 38.51 (265.5)

MECHANICAL PROPERTIES DATA CENTER

2 Q4

Table 3. Tabular presentation of the best-value curve for the AISI 316

data shown in Fig. 1.

513
PERCENT

MATERIALS FOR

SUPERCONDUCTING

MAGNET SYSTEMS

AlSl 316 Structural Alloys

Stainless Steels

0 CO 120 180 240 300 TEMP - K

THE ELONGATION OF ANNEALED TYPE 316, 316L, 316LN

STAINLESS STEEL (1, 2, 3, 4, 6, 7, 9, 10, 11, 12)

0 60 120 160 240 300 TEMP - K

THE REDUCTION IN AREA OF ANNEALED TYPE 316, AND 316LN STAINLESS STEEL (1, 2, 3, 4, 6, 9, 10, 11)

MECHANICAL PROPERTIES DATA CENTER

Figure 2. A typical handbook graph - B-category property. 2.05

514
VAIL WORKSHOP

VAIL WORKSHOP

F. R. Fickett

Thermophysical Properies Division

National Bureau of Standards

A major aspect of the technology transfer part of the program is the

workshop. These meetings are held each year with the goal of bringing

together representatives from all parts of the materials community who

are involved with low temperature applications. The attendees at this

year's workshop, held at Vail, CO, October 24-26, represented just such

a diverse group including: industry, national labs, MFE projects and

universities. The fol 1 owing -pages give the attendance list and the program

for the workshop. No proceedings are published, but extended abstracts of

the presentations were provided to the attendees. In addition a review of

the workshop will be presented at the First Topical Meeting on Fusion Reactor Materials, January 29-31, 1979.

317
PROGRAM FOR NBS-DOE WORKSHOP ON

MATERIALS AT LOW TEMPERATURES

Vailj Colorado October 24-26, 1978

Tuesday, Oct. 24

I. INTRODUCTION (8:30-9:30), E. N. C. Dalder, Chairman

Welcome F. R. Fickett NBS

OFE overview E. N. C. Dalder DOE

NBS program R. P. Reed NBS

Handbook project T. Moore/B. Gilmore MPDC

II. DESIGN I (10:00-12:00), F. R. Fickett, Chairman

Large coil project C. J. Long

E. Belkin/T. Leax

F. Mazandarany

J. Christian

ORNL

Westinghouse

General Electri

General Dynamic

III. DESIGN II (2:00-4:30), F. R. Fickett, Chairman

MFTF J. Hodges LLL

MHD magnets R. Shaw MIT

Pulsed magnets Z. N. Sanjana

S. T. Wang

Westinghouse

ANL

TFTR

Wednesday, Oct. 25

G. Brown PPPL

IV. ALLOYS FOR LOW TEMPERATURE SERVICE (9:00-12:00), R. P. Reed, Chairman

Fatigue J. Shepic Martin Marietta

Fracture & tensile D. T. Read NBS

Production H. Espy Armco

Elastic properties H. M. Ledbetter NBS

Magnetic properties T. Col lings

F. Fickett

Battel le

NBS

V. THICK SECTION WELDING FOR LOW TEMPERATURE APPLICATIONS (2:00-5:00), H. I. McHenry, Chmn

Stainless steel weld metallurgy

Cryogenic toughness of stainless

steel weld metals

Influence of delta ferrite on toughness

Nickel -based weld metals

NBS/DOE welding program

01 son Reid

D. T. Read

C.

H.

Wi therel

1

I. McHenry

Colo. School of Mines

Teledyne-McKay

NBS

LLL

NBS

EVENING (6:00-8:00) Fluid flow studies in the lobby bar area (cocktail party)

Thursday, Oct. 26

VI. NONMETALLICS FOR MAGNET STRUCTURES (8:30-12:00), M. B. Kasen, Chairman

ICMC Munich Conference review

NBS/DOE composites program

Composite production

Radiation damage

Composites testing

Elastic properties

Friction and wear of polymers

A.

M.

J.

C.

E.H.

R.

F. Clark

B. Kasen

Benzinger

J. Long

Erez

Ledbetter

Tobler

NBS

NBS

Spaulding

ORNL

MIT

NBS

NBS

Fiber
ATTENDANCE

AT

NBS/DOE

VAIL

WORKSHOP

October

24-26,

1978

i/> as-fO ac

fO

-(-> t/) M-O

3

fOa»

CO

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03 03 X3 O 03

03 Q

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</3 4-O

CO03

03

CO -a

s_03

• -o

JZl C

03 03

CO "O

S-03 X3e

03

i- X) 4-3 1 u 03 0 <>Q c>o E CJ 03 03 •r- E 4- E >1 CQ _J 03 E r— 4-3 03 0 E 03 4- 4- f— O UJ E cu 03 r— 0 0 r— UJ 03 E > 03

03 ra 4-3 CL 3 0 03 •r- -X 13 C3 E E 03 rU O CQ 0 E —1 E 03 03 O o XI 03 2 E 2 E 03 03 03

•r- 03 E 03 Q, r— u 0 03 CQ E E

03 4-3 4-3 f— 03 03 E 03 03 c/3 CT3 U C3 Z3

() fO fO E CD 03 n— E E 1— E E 03 CQ OQ

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NBS-114A (REV. 9-78)

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBSIR 79-1609

ltec«$»ion No. Z, ftociptent’s Acco$(aion No,

4. TITLE AND SUBTITLE 5. Publication Date June 1979

Materials Studies for Magnetic Fusion Energy Applications

at Low Temperatures - II

8. Performing Organization Code

7. AUTHOR(S)

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

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, DC 20234

Ift, PR>ject/Task/Woid( tbiit No.

'' 7360531 ‘ :

11. Contract/Grant No.

12. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS rsrreer, City, state, ZIP)

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

13. Type of Report & Period Covered

15. SUPPLEMENTARY NOTES

I I Document describes a computer program; SF-185, FIPS Software Summary, is attached.

16. ABSTRACT (A 200^word or less factual summary of most significant information. If document includes a significant bibliography or

literature survey, mention it here.) The reports presented here summarize the work of the low temperature materials

research project for the second year of the program. The various projects are outlined and the research results are presented. The major thrust of the measurements

has been the evaluation of the low temperature mechanical and physical properties of

stainless steel base metal and welds, with particular emphasis on the nitrogen

strengthened stainless steels. Aluminum alloys have received some consideration also

Work has been done on the production of nonmetallics, primarily industrial laminates

for low temperature applications and on the measurement of their properties at cryogenic temperatures. The second NBS/DOE Vail workshop was held in October 1978. A

brief description is given of that program.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name;

separated by semicolons)

Alloys; fiberglass; mechanical properties; nonmetallics; physical properties;

stainless steel.

18. AVAILABILITY [^Unlimited

I I For Official Distribution. Do Not Release to NTIS

I I Order From Sup. of Doc., U.S. Government Printing Office, Washington, DC

20402, SD Stock No. SN003-003-

rX~l Order From National Technical Information Service (NTIS), Springfield, VA. 22161

19. SECURITY CLASS

(THIS REPORT)

UNCLASSIFIED

20. SECURITY CLASS

(THIS PAGE)

UNCLASSIFIED

21. NO. OF

PRINTED PAGES

496

22. Price

$12.50

ir U.S. GOVERNMENT PRINTING OFFICE;1979—677-072 / 1291 USCOMM-DC

V