NATIONAL INSTITUTE OF STANDARDS &

TECHNOLOGY

Research Information Center

Gaithersburg, MD 20899

A11100 T1077k

NBSIR 81-1645

NESS

UCATIONS

TECHNICAL REPORTS

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES - IV

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

By

Fracture and Deformation Division

L National Bureau of Standards

^ Boulder, CO 80303

j56

31-1645

1981

:.2

NBSIR 81-1645

«-* ¥ TECHNICAL REPORTS

LIBRARY

JUt. 2 0 1981

MATERIALS STUDIES FOR

MAGNETIC FUSION ENERGY

APPLICATIONS AT LOW TEMPERATURES - IV

Edited By

R.P. Reed and N.J. Simon

Fracture and Deformation Division

National Bureau of Standards

Boulder, CO 80303

April 1981

Sponsored By

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

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

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

CONTENTS

Page

SUMMARY v

PROGRAM DESCRIPTION vii

HIGHLIGHTS OF RESULTS 1

STRUCTURAL ALLOYS

STRUCTURAL ALLOYS PROGRAM 11

STRUCTURAL ALLOYS FOR SUPERCONDUCTING MAGNETS IN

FUSION ENERGY SYSTEMS 17

STRENGTH AND TOUGHNESS RELATIONSHIP FOR INTERSTITIALLY

STRENGTHENED AISI 304 STAINLESS STEELS AT 4 K 37

TENSILE AND FRACTURE PROPERTIES OF MANGANESE-MODIFIED

Fe-18Cr-8Ni TYPE AUSTENITIC STAINLESS STEELS 77

INTERSTITIAL CARBON AND NITROGEN EFFECTS ON THE CRYOGENIC

FATIGUE CRACK GROWTH OF AISI 304 TYPE STAINLESS

STEELS 101

ANOMALOUS YIELDING, TRANSFORMATION, AND RECOVERY EFFECTS

IN AISI 304L STAINLESS STEELS 131

LOW-TEMPERATURE DEPENDENCE OF YIELDING IN AISI 316

STAINLESS STEELS 147

LOW-TEMPERATURE DEFORMATION OF 5083-0 ALUMINUM ALLOY 185

EFFECTS OF CARBON AND NITROGEN ON THE ELASTIC CONSTANTS

OF STAINLESS STEEL 304 203

LOW-TEMPERATURE VARIABILITY OF STAINLESS-STEEL-304

ELASTIC CONSTANTS 215

PREDICTED SINGLE-CRYSTAL ELASTIC CONSTANTS OF STAINLESS

STEEL 304 227

ELASTIC-CONSTANT/TEMPERATURE BEHAVIOR OF THREE HARDENED

MARAGING STEELS 237

TEMPERATURE BEHAVIOR OF YOUNG'S MODULI OF FORTY ENGINEERING ALLOYS 257

WELDING AND CASTING

EVALUATION OF WELDMENTS AND CASTINGS FOR LIQUID HELIUM

SERVICE 273

STRENGTH AND TOUGHNESS OF FULLY AUSTENITIC STAINLESS

STEEL FILLER METALS AT 4 K 289

WELD PROCESS STUDY FOR 316 L STAINLESS STEEL WELD METAL

FOR LIQUID HELIUM SERVICE 303

FRACTURE TOUGHNESS PROPERTIES OF WELDS IN ALUMINUM ALLOY

5083-0 AT 4 K 323

THE INFLUENCE OF FERRITE AND NITROGEN ON MECHANICAL PROPERTIES OF CF8M CAST STAINLESS STEEL 337

METALLOGRAPHY OF 6-FERRITE IN AUSTENITIC STAINLESS STEEL

CASTINGS 357

DEFORMATION AND FRACTURE OF STAINLESS STEEL CASTINGS AND

WELDMENTS AT 4 K 415

EVALUATION OF WELD FILLERS TO JOIN NITROGEN STRENGHTHENED

AUSTENITIC STAINLESS STEELS FOR CRYOGENIC APPLICATIONS (FUSION AND MHD ENERGY) 453
NONMETALLICS

NONMETALLICS FOR MAGNET SYSTEMS 467

NONMETALLIC MATERIALS: STANDARDIZATION

STANDARDIZING NONMETALLIC COMPOSITE MATERIALS FOR

CRYOGENIC APPLICATIONS 477

NONMETALLIC MATERIALS: RESEARCH AND CHARACTERIZATION

THERMAL CONDUCTIVITY OF G-IOCR AND G-11CR INDUSTRIAL

LAMINATES AT LOW TEMPERATURES (PART III) 495

FRACTURE ANALYSIS OF COMPRESSION DAMAGE IN GLASS FABRIC/

EPOXY, G-IOCR, 295 K to 4 K 507

STATIC AND DYNAMIC YOUNG'S MODULI OF TWO FIBER-REINFORCED

COMPOSITES 589

RADIATION EFFECTS ON COMPOSITES AND NONMETALLIC

MATERIALS 597

NONMETALLIC MATERIALS: TECHNOLOGY TRANSFER

CONTRIBUTIONS TO TECHNOLOGY FORECAST '81 605

COMMENT ON "MECHANICAL PROPERTIES OF PLASTIC COMPOSITES.." 609

THE MANUFACTURE OF RADIATION RESISTANT LAMINATES 613

TECHNOLOGY TRANSFER

NBS-DoE WORKSHOP: MATERIALS AT LOW TEMPERATURES 641

ORGANIZATIONAL CONTACTS 646

ACKNOWLEDGMENT 647
SUMMARY

This report contains results of a research program to produce

material property data that will facilitate design and development of

cryogenic structures for the superconducting magnets of magnetic fusion

energy power plants and prototypes. The program was conceived and

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

the Office of Fusion Energy of the Department of Energy; it is managed

by NBS and sponsored by DoE. Research is conducted at NBS and at

various other laboratories through subcontracts with NBS.

The reports presented here summarize the fourth year of work on the

low temperature materials research program. Highlights of the results

are presented first. Research results are given for the four main

program areas: structural alloys, weldings and castings, nonmetal 1 ics,

and technology transfer. Objectives, approaches, and achievements are

summarized in an introduction to each program area.

The major portion of the program has been the evaluation of the low

temperature mechanical and physical properties of stainless steel base

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 and standardization of

nonmetallics, primarily industrial laminates, for low temperature applications and on the measurement of their properties at cryogenic

temperatures . A brief description is given of the fourth NBS/DoE Vail

workshop held in October 1980.

v
Note

Certain commercial equipment, instruments, or materials are

identified in this paper to specify the experimental procedure adequately. In no case does such identification imply recommendation or

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

the material or equipment identified is necessarily the best available

for the purpose. Papers by non-NBS authors have not been reviewed or

edited by NBS . Therefore, the National Bureau of Standards accepts no

responsibility for comments or recommendations contained therein.

vi
PROGRAM DESCRIPTION

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

construction, and safe operation of low temperature magnetic fusion

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

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

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

personal contacts, publications, and workshops.

Efforts directed at the first specific objective began with the

publication of the Survey of Low Temperature Materials for Magnetic

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

survey, entitled Structural Alloys for Superconducting Magnets in Fusion

Energy Systems , is included in this volume. Through continuous interactions

with all low temperature design, construction, and measurement programs,

such as the Large Coil Project, we are aware of new problems as they

arise. This year's contribution to the second objective is described in

Table 1 in the form of an outline of the research projects. The results

appear later in this report. The third objective is satisfied, in part,

by these annual reports and by the series of NBS-DoE Workshops on Materials at Low Temperatures, which are held each fall in Vail, Colorado.

Starting in fiscal year 1981, handbook pages presenting the available

data for specific materials will be added to the annual report.
Table 1. Outline of the NBS/DoE program on material studies for

magnetic fusion energy applications at low temperatures.

Program Area

1 . Characteri zation

Organ i zation

A. Structural Alloys

NBS

2. Alloy development NBS

3. Characterization and U. Wisconsin

modeling of martensitic

start temperatures

B. Welds and Castings

1. Stainless steel welds NBS

2. Aluminum welds NBS

Program Description

Measurement of elastic, tensile,

fracture toughness and fatigue

crack growth rate properties as

a function of temperature for

selected austenitic stainless

steels and 5083 aluminum. Assessment of effects of grain size,

martensitic transformations and

alloying on low temperature behavior .

Assessment of low temperature

properties of selected austenitic alloys containing varying

concentrations of C, N and Mn

with a Fe-Cr-Ni base.

Study to assess, empirically,

the relation between alloy content and the temperature at

which martensite forms on cooling

for commercial grades of AISI 300

series stainless steels.

Evaluate the strength and

toughness of several fully aus- tenitic weld metals provided by

suppliers in the United States,

Austria, Sweden, and the Soviet

Union at 4 K. (Four alloys with

16-20% Cr, 16-20% Ni , 4-7% Mn

and 2-5% refractory elements (Mo,

W and/or Nb) had strength and

toughness combinations comparable

to base metal properties.)

Evaluate the strength and fracture

toughness of GMA welds using 5183

wire in 51 mm thick 5083-0 plates

that were provided by five aluminum fabricators. (Property

variations were minimal despite

significant differences in

welding procedures.)

v i i i
3. Stainless steel

castings

1. Standardization

3. Characterization

1 . NBS/DoE Workshop

Produce stainless steel castings

(grade CF8M) with varied ferrite

and nitrogen contents to study

the effects of ferrite and nitrogen on the strength and toughness at 4 K. (This year's efforts consisted of manufacture

of the castings, tests at room temperature and 77 K, and metal lographic studies to be followed by

evaluation in liquid helium next

year.

)

C. Nonmetallics for Magnet Structures

NBS Continue cooperation with industry to establish and refine

cryogenic grade specifications

for insulating laminates and

coding systems for laminates.

NBS Measure thermal conductivity of

composite components at cryogenic

temperatures to aid in establishing models for predicting the

thermal conductivity of a variety

of laminates. Investigate failure

modes of laminates under compression at cryogenic temperatures.

Develop dynamic methods of measuring elastic constants of com- posite laminates.

ESCO

NBS

D. Technology Transfer

NBS An annual workshop to present re- search results to the fusion com- munity, to discuss new problems,

and to promote interaction between

interested parties. Handbook

data pages on stainless steels

will be produced, starting in

FY '81.

ix
wmm
HIGHLIGHTS Of RESOL TS

1

HIGHLIGHTS OF RESULTS

STRUCTURAL ALLOYS:

Stainless Steel 3Q4--Characterization and Alloy Development

The study of this alloy series has been emphasized because of its

favorable cryogenic properties, extensive service record, and availability. Stainless 304 has been used in several designs for the Large

Coil Project (LCP). Previous work has shown that adding nitrogen

increases the tensile strength and austenitic stability of the 304

series at low temperatures. Carbon has a similar effect, but the addition of carbon is limited owing to sensitization (chromium carbide

precipitation). It would be desirable to optimize the properties of the

alloy for different requirements by varying the composition, and this

objective has motivated the studies described below.

1. Carbon and nitrogen (C+N) weight percent were varied systematically in a series of heats for which fracture toughness and

tensile yield strength were measured at 4 K. Fracture toughness dropped linearly with increasing C+N content, but strength

increased as C+N content increased.

2. The effects of C+N content and temperature upon fatigue crack

growth were also measured. Low C+N content was found to

correlate with better crack resistance at 4 K. The improved

crack resistance for low C+N contents at 4 K, compared with

that at room temperature was shown to be due to a transition

to transgranular fracture at cryogenic temperatures from

intergranular at room temperature.

3
3 . Although interstitial nitrogen strengthens the 304 alloy

series, the resulting alloys are harder to fabricate. Since

the addition of manganese is expected to improve fabricabil ity,

a series of heats with manganese content varying from 1 to 6

weight percent and nitrogen content of 0.1 to 0.2 weight

percent were tested for fracture toughness and tensile yield

strength. The results gave much lower fracture toughness than

expected, except for the heat with 6 percent manganese. These

anomalous results may indicate the presence of previously

unrecognized factors in alloy production that affect fracture

toughness significantly at cryogenic temperatures. Research

to elucidate these factors is continuing.

4. Yield strength of a selected heat of 3041 was measured as a

function of temperature. The measurements showed that yield

strength does not increase monotonical ly with decreasing

temperature, as is the case for the austeni tical ly stable

stainless 310.

5. Effects of C+N content on the elastic constants of 304 steels

were measured at room temperature, and the variability among

the elastic constants of alloys of the 304 series was measured

down to 4 K.

Stainless Steel 316

The low temperature flow strength of stainless 316 was measured.

The formation of bcc martensite that occurs at low temperatures or as

the result of strain influences this property.

4
Aluminum Alloys

The yield strength and stress-strain curves of the aluminum alloy

5083-0 were measured down to 4 K. Serrated yielding was observed at 4 K

and room temperature, though not at 76 K. The mechanism causing this

effect is different at 4 K than at room temperatures

.

The magnesium present in this alloy increases the strength at both

room and cryogenic temperatures. Previously unpublished data on alloys

5154 and 5052, which have different magnesium content, are presented to

illustrate this trend further.

WELDING AND CASTING:

Research has focused on the metallurgical factors that determine

strength and toughness of weldments and castings in stainless steel and

aluminum alloys.

1. Since it is known that ferrite reduces the fracture toughness

of weldments at 4 K, although ferrite reduces microfissuring

,

work on the development of 0 percent ferrite filler materials

has been undertaken. A series of experimental filler alloys

was tested at 4 K. Some were found to have a good combination

of strength and toughness, comparable to that of the base

metal, but a few problems with microfissuring remain.

2. The effects of differing weld processes were studied for 31 6L

weldments. The cleanliness of the process correlated well

with good fracture toughness at 4 K, so it was concluded that

the impurity level may have a strong effect on cryogenic

properties of welds.

5
3 . To support LCP magnet design and construction, a number of

qualifying and test weldments of stainless steel were tested

for fracture toughness and tensile strength.

4. The evaluation of 5183 welds in 5083 aluminum alloys was

completed during the past fiscal year. The fracture toughness

of weldments at 4 K was found to be essentially the same for

welds furnished by different suppliers or produced by different

processes

5. Work has begun on a study of the more fundamental aspects of

the structure of duplex austeni te/ferri te steels and on how

this structure is related to fracture and deformation in weldments and castings at cryogenic temperatures. Castings are

being studied for two reasons: (a) there is interest in their

use as structural elements for superconducting magnets and

(b) the factors affecting fracture and deformation modes are

more easily observed because their ferrite structure is coarser

than that of weldments.

A series of CF8M castings with varied ferrite and nitrogen content

was produced this year, and the production techniques are reported.

This series will be used to study the effects of ferrite and nitrogen in

a controlled manner in the current fiscal year. Meanwhile, during the

past year, available weldments and castings were subjected to a detailed

metal lographic study. Fracture and deformation modes were determined

for these specimens, and the effect of ferrite structure on these modes

was analyzed.

6
NONMETALLICS

:

Standardization

Efforts to refine and improve the specifications of cryogenic

grades of glass-fabric-reinforced bisphenol A epoxy industrial laminates

have continued. These established grades are known as G-10CR and G-11CR.

Recently it was found that G-11CR does not have the radiation resistance

that was initially expected from its composition and curing process.

Therefore, this grade has been dropped. The standardization of glassmat reinforced laminate products is in progress.

Characterization

The thermal conductivity of the epoxy matrix material used in

G-10CR and G-11CR was measured, and the results revealed considerable

inhomogeneity in the bulk resin. Efforts to compare thermal conductivity measurements made at the National Bureau of Standards with those

made by CERN on a laminate of European manufacture indicated a sizeable

discrepancy at room temperature. Work on this problem will be continuing during the present fiscal year.

Modes of failure in compressive fracture of laminates at cryogenic

temperatures were studied. Compressive stresses of insulators predominate in superconducting magnets.

Dynamic elastic moduli were measured for two fi ber-rei nforced composites, and the results were compared with static methods.

Results of studies on radiation effects on composites are presented

in a section from a review paper.

TECHNOLOGY TRANSFER:

The annual NBS/DoE Workshop on Materials at Low Temperatures at

Vail, Colorado in October 1980 was attended by about seventy representatives

7
of the materials and magnet design communities. Again, the meeting was

viewed constructively ; it served as a focal point to identify and discuss magnet construction or design problems related to materials. For

a few who have become recently involved in cryogenic applications, the

workshop served an educational role.

8
STRUCTURAL ALLOYS

STRUCTURAL ALLOYS PROGRAM

R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

An overview of the structural alloys program is presented. Selection and properties of the alloys under study are discussed, and the

objectives of the program in the current fiscal year (1981) are set

forth.

11

The objectives of the structural alloys project are twofold:

(1) to characterize the physical and mechanical properties of structural

alloys at 4 K to obtain typical properties than can be used for material

selection and design estimates and (2) to develop, in collaboration with

industry, a stronger, tough, weldable austenitic steel for use as a

structural alloy for superconducting magnets.

The structural -al loys-project results from this past year are summarized :

Characterization : The tensile yield strength as a function of

temperature (4-330 K) was measured for a thick

plate of 5083-0 aluminum alloy.

The temperature dependence of the flow strength

of AISI 316 was measured and the influences of

martensitic transformations, adiabatic heating,

magnetic transitions, and grain size were

assessed

.

Measurements on fatigue and creep of metastable

AISI 304 were initiated.

The fatigue crack growth rate of a series of

Fe-18Cr-8Ni alloys with varying C and N concentrations was measured and assessed.

Measurements of elastic constants of stainless

and maraging steels, of the variability of

AISI 304 elastic properties and a review of the

temperature dependence of elastic constants of

engineering alloys are presented.

13
Alloy Development : The tensile yield and ultimate strengths and

the fracture toughness of a series of Fe-18Cr7Ni-0.03C alloys with varying Mn and N concentrations were measured at 295, 76, and 4 K.

Fracture toughness results obtained for the low

(1-4 wt. %) Mn alloys were anomalous when

compared with the Fe-1 8Cr-8Ni-(C+N) alloy

series measured the previous year.

This year (fiscal year 1981), the tensile flow strength of 2219-T87

aluminum alloy will be measured at a series of temperatures ranging from

4 to 300 K. Additional tensile measurements will be conducted on 5083-0

aluminum alloy to assess the influence of grain size on strength and to

assess, more accurately, the temperature dependence of the flow strength.

The study of the role of martensitic transformation on low- temperature

fatigue and creep of AISI 304 steel will continue. Finally, to pursue

the development of a stronger, but tough, austenitic steel, the causes

for the lower toughness at 4 K of Fe -18Cr - 7Ni -l-4Mn-N steels of the

last alloy series need to be understood. Metal lography , electron microscopy, scanning electron microscopy, and scanning transmission

electron microscopy will be used to study the fracture surfaces, crack

development, and impurity segregation of these alloys. Subsequently, an

alloy heat of selected composition will be prepared.

The program has concentrated on four alloy series: AISI 304 (Fe1 8Cr-8Ni ) , AISI 316 ( Fe-1 6Cr-l ONi -2Mo) , 5083-0 ( A1 -0. 7Mn-4. 45Mg-0 . 01 5Cr)

and 2219 (Al-6Cu-0.3Mn-0.lV-0.18Zr-0.06Ti). These alloys are most

commonly used for 4 K applications and are thought to represent the best

base alloy materials.

14
The stainless steels have a relatively high Young's modulus, low

thermal and electrical conductivity, and excellent toughness. Also, the

major fabricators have good experience in welding stainless steel.

Aluminum alloys are light, less expensive, and have been used extensively (alloy 5083) in the construction of LNG containment systems.

Toughness and strength do not decrease between 110 K and 4 K for these

aluminum alloys.

AISI 304 is susceptible to martensitic transformation on cooling or

on application of applied stress at low temperatures . Martensitic

transformation tends to act as a deformation mode, therefore its presence tends to reduce the resistance of the alloy to applied stress at

low temperatures , and the result is low initial flow strengths. However, we have not obtained evidence that the martensitic transformations

act to degrade fracture toughness or ultimate tensile strength.

The alloy AISI 316 is less susceptible to martensitic transformation and the dependence of strength on temperature for this steel is

more predictable.

The addition of nitrogen to austenitic stainless steel dramatically

increases the 1 ow-temperature yield strength of the steel. There is an

attendant decrease of toughness. The aim of the alloy development

program is to establish a composition of Fe-Cr-Ni-Mn-C-N that will

provide a stronger and tough structural alloy at 4 K. Two correlations

have been established to assist this process: (1) At 4 K the tensile

yield strength has been demonstrated to be inversely related to the

fracture toughness ( J-integral ) . (2) The total concentration of C+N

from about 0.10 to 0.30 wt. t is linearly related to the tensile yield

strength at 4 K.

15
The aluminum alloys do not exhibit low-temperature phase or magnetic transitions. Consequently, both their physical and mechanical

properties are well behaved at very low temperatures. The fracture

toughness of the aluminum alloys is considerably lower than stainless

steel values. However, there is a range of slow, stable crack growth

following crack initiation at the critical J-integral fracture toughness

and additional research is required to characterize the range of stable

crack growth.

16
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

17

Nuclear Engineering and Design 58 (1980) 219-236

© North-Holland Publishing Company

STRUCTURAL ALLOYS FOR SUPERCONDUCTING MAGNETS IN FUSION ENERGY SYSTEMS

H.I. McHENRY and R.P. REED

Thermophysical Properties Division, National Bureau ofStandards, Boulder, Colorado 80302, USA

Received 15 March 1979

The behaviour 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), nitrogen-strengthened austenitic

stainless steels (types 304LN, 316 LN and 21Cr-6Ni-9Mn) and aluminum alloys (grades 5083, 6061 and 2219). The

mechanical and physical properties of the selected alloys at 4 K are reviewed. Welding, the properties of weldments, and

other fabrication considerations are briefly discussed. The available information suggests that several commercial alloys have adequate properties at 4 K and sufficient fabrication characteristics for the large magnet structures needed for fusion energy systems.

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 accessibility 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 fabricability to permit manufacture and assembly of the components.

The choice of materials is further restricted by the adverse operating environment of the magnet system.

The alloys are subjected to temperatures down to

4 K and must not be susceptible to low temperature

embrittlement. The presence of high magnetic fields of controlled shape cautions against the use of ferro- magnetic 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 coils cause heat

inputs to the structure from eddy current losses. Consequently, alloys with high electrical resistivity are usually preferred. The radiation environment

eliminates use of alloying elements, such as cobalt,

which transmute to long half-life products. With this exception, radiation does not appear to be a signifi- cant factor in the selection of structural alloys

because dosage levels sufficient to degrade structural performance probably cause unacceptable deter- ioration 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, ferro- magnetic phase and the attendant volume expansion

causes localized areas of stress concentration.

The contraints on materials selection caused by

size requirements, space limitations and operating

environment significantly reduce the number of can- didate alloys. Further limitations arise due to con19
//./. McHenry, R.P. Reed / Structural alloys for superconducting magnets

siderations 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 4 K 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 US Department of Energy. For background

relating to the properties discussed the reader is refer- red 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 4 K. The

physical properties also offer advantages over competing materials. Specifically, the modulus is high, the thermal expansion is close to that of the copperstabilized 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 microstructural

stability and control. The AISI 300-series stainless steels, particularly

types 304 and 304L, are widely used in cryogenic

applications. The designations and compositions of

commonly used grades are summarized in table 1. These

grades have moderate strength, excellent thoughness,

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 coverTable 1 Compositions of austenitic stainless steels suitable for liquid helium service

AISI Composition (%)

Cr Ni C, max Other

304 18-20 8-10.5 0.08 304L 18-20 8-12 0.03 316 16-18 10-14 0.08 2-3 Mo

316L 16-18 19-22 0.03 2-3 Mo

310 24-26 19-22 0.25 1.5 Si, max

310S 24-26 19-22 0.08 1.5 Si, max

For each grade Mn = 2.0 max; S = 0.030 max; Si = 1.0 max;

P = 0.045 max.

ed by ASTM specifications and can meet ASME Pres- sure Vesels Code requirements for cryogenic vessels. These grades are not described in this chapter because

they offer no particular advantage for superconducting magnet structures and consequently their service experience and property data at 4 K are limited. The alloys listed in table 1 can be compared with

respect to AISI 304, the basic 19Cr—9Ni 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 car- bon is desirable to avoid sensitization, the grain boundary precipitation of Cr-carbides. Type 31 OS is a 25Cr—20Ni alloy that has lower carbon (0.08% C)

than AISI 310 (0.25% C). The high alloy content

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 A 194

20
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

provides austenite stability and greater strength and

thus 31 OS is useful where dimensional stability, non

magnetic behavior and higher strength are desirable. The low temperature properties of the 300-series

austenitic 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—0.08 wt% is particularly significant with the higher

nitrogen contents causing substantial increases in yield

strength.

2.1. Tensile properties

The yield and ultimate strengths of AISI grades

304, 310 and 316 are compared in fig. 1 for temperatures ranging from 4 to 300 K. 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 fig. 2. Type 304

exhibits a slightly greater increase in strength than type

304L as temperature is reduced. The loss of strength

associated with reduced carbon content is considered

to be characteristic of all the stainless steels, i.e. 304L

vs. 304, 316L vs. 316 and 31 OS vs. 310, and is attri- buted to the strengthening effect of carbon.

The ductility of the annealed 300-series stainless

steels is generally exce'lent at cryogenic temperatures.

Elongation and reduction-of-area tend to drop with

decreasing temperatures, but values generally exceed

30%.

Fig. 1. Tensile and yield strengths of three austenitic stainless

steels - AISI types 304, 310 and 316 - at temperatures between 4 and 300 K [6].

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 0.38 mm lateral expansion requirement

of the ASME Boiler and Pressure Vessel Code [11]

and a ductile-to-brittle transition is not exhibited. The

minimum temperature for Charpy testing is about

20 K, because adiabatic heating causes specimen temperatures in excess of 20 K even when the test temperature is 4 K. Slow strain-rate tests using notched tensile speci- mens indicate that notch sensitivity does not develop

at 4 K [6] . The data summarized in fig. 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.

21
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Fig. 2. Tensile properties of two austenitic stainless steels —

AISI types 304 and 304L — at temperatures between 4 and 300 K [6],

Fig. 3. Notched tensile properties of five austenitic stainless

steels - AISI types 304, 304L, 310, 310S and 316 - at 4 K

[6]. K-p is the stress concentration factor of the notch.

Tobler [12] used /-integral methods to measure tht fracture toughness of AISI grades 310 and 316 at temperatures to 4 K. The toughness, K\C(J), at 4 K is 230

MPaVm for 31 OS and 460 MPa\/m for 316, and in both alloys the toughness at 4 K exceeds the toughness at room temperature. The ratio of toughness to

yield strength is sufficiently high to ensure gross ductile deformation prior to fracture.

2.3. Fatigue

Strain cycling fatigue properties at 300, 76 and

4 K have been measured for AISI grades 304L and

3 1 0 by Nachtigall [13] and for grades 304L and 3 1

by Shepic and Schwartzberg [14], The results for 304L, which were essentially the same in both investi- gations, and for 316 are shown in fig. 4. The fatigue

resistance of 310 and 316 were superior to that of

304L, particularly in low cycle fatigue (less than 104

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 4 K by Tobler and Reed [15]. The data for 304L are shown in fig. 5. Notice that the growth

rates are essentially the same at 76 and 4 K; similar behavior was observed in each of the other alloys.

Also, the growth rates at 76 and 4 K are slower than

those at room temperature. The best-fit lines through

the 304L data at 295 K and at 76 and 4 K form the

Fig. 4. Fatigue behavior of three austenitic stainless steels — AISI types 304 and 316 and a 21Cr-6Ni-9Mn alloy — at 4 K [ 14]

22
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Fig. 5. Faitgue crack growth rates of AISI type 304L austenitic stainless steel at 295, 76 and 4 K (15].

approximate scatter bands for the data on 304, 31 OS

and 316 at 295, 76 and 4 K. Thus, it can be concluded

that alloy content and temperature have minimal

influence on the fatigue crack growth behavior of

the 300-series stainless steels.

2.4. Physical properties

The physical properties of the 300-series stainless

steels at 295, 76 and 4 K 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 304 N (AISI 304 with deliberate additions

of nitrogen) is approximately three times greater than

the yield strength of 304 at 76 K, the ultimate strengths

are approximately equal and the elongation of the 304N is superior at temperatures below 220 K [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

Table 3 Physical Propeties of austenitic stainless steels

Property Units AISI 304 AISI 310 AISI 316

295 K 4 K Ref. 295K 4 K Ref. 295 K 4 K Ref.

Density g/cm 3 7.9 _ 19 7.9 _ 19 8.0 19 Young’s modulus GN/m2 189 204 16 184 199 16 217 234 16 Shear modulus GN/m2 740 800 16 700 770 16 830 910 16 Poisson’s ratio - 0.290 0.271 16 0.307 0.296 16 0.298 0.281 16

Thermal conductivity W/m • K 13 0.3 7 nearly the same for 304, j310 and 316

Expansion cm/cm • K X 10~6

15.8 10.2 7 nearly the same for 304, 310 and 316

Specific heat J/kg • K 480 0.5 7 nearly the same for 304, 310 and 3 16

Electrical resistivity p Dcm 70.4 49.6 17 87.3 68.5 17 75.0 53.9 17 Magnetic permeability 1.02 1.09 18 1.003 1.10 18 1.003 1.02 18

23
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Table 4 Compositions of nitrogen-strengthened austenitic stainless steels suitable for liquid helium service

Type Nominal composition (%)

Cr Ni Mn C, max N Other

AISI 304N 18 -20 8 -10.5 2.0, max 0.08 0.10-0.16 _

304LN 18 -20 8 -12 2.0, max 0.03 0.10-0.16 -

AISI 316N 16 -18 10 -14 2.0, max 0.08 0.10-0.16 -

316LN 16 -18 10 -14 2.0, max 0.03 0.10-0.16 - ASTM XM1 1 * 18 -21 5 - 7 8-10 0.04 0.15-0.40 - ASTM XM19 * 20.5 -23.5 11.5 -13.5 4- 6 0.06 0.20-0.40 Cb 0.10-0.30 V 0.10-0.30 Mol. 5 -3.0

* ASTM A41 2.

nitrogen is deliberately added to several of the com- mon Cr—Ni 300-series stainless steels. The equivalent AISI designations for these European grades are 304 LN and 316LN.

Two basic types of nitrogen-strengthened stainless

Fig. 6. Tensile and yield strengths of nitrogen-strengthened

300-series stainless steels at temperatures between 4 and 300 K [20-27).

steels are suitable for cyrogenic service: the Cr—Ni—

and the Cr-Ni-Mn-N alloys. The Cr-Ni-N alloys

are 300-series stainless steels with deliberate additions of nitrogen. In the US specifications, the nitrogen

range is 0.10—0.16% for 304N and 316N. The corres- ponding European specifications permit nitrogen

levels to 0.25%. The manganese in Cr—Ni—Mn—

alloys 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—

300-series stainless steels are summarized in fig. 7. The solid lines and the broken lines represent the data of Randak et al. [21] for 304LN and 316LN,

respectively. 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 in the case for 304 and 3 1 6, fig. 1 , alloys 304N

and 304LN have approximately 20% greater ultimate

strength at cryogenic temperatures than alloy 316LN.

All three alloys have essentially the same yield

strength at and below room temperature. 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 stain24
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Fig. 7. Tensile and yield strengths of 21Cr—6Ni-9Mn stainless steel at temperatures between 4 and 300 K [ 28].

less 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

Fig. 6. This alloy typically has about 25% greater

yield strength than the Cr—Ni—N alloys. Representative tensile data on the other nitrogen-strengthened

Cr—Ni—Mn alloys that have been tested at 4 K indicate

that yield strength increases and ductility decreases with decreasing temperature; and thus, fracture resis- tance must be considered when selecting these alloys

for liquid helium service.

3.2. Toughness

The nitrogen-strengthened stainless steels have

excellent toughness at room temperature. However,

as shown in fig. 8, significant toughness losses generally occur as the temperature is reduced [25]. Read

and Reed [25] observed significant variations in the toughness of a single piece of 21 Cr— 12Ni-5Mn alloy.

As-received mill-annealed material having an inter- granular microconstituent had a plane strain fracture toughness at 4 K of 1 1 1 MPa\/m. A high temperature

(1177°C, 1.5 h) anneal dissolved the microconstituent

and increased the toughness to 176 MPa\/m- Thus,

certain nitrogen-strengthened grades can be used at 4 K, but care should be taken to ensure that the alloy selected has satisfactory toughness at 4 K for the

applicable melting practice, product form and heat

treatment.

TEMPERATURE, K

Fig. 8. Fracture toughness of four nitrogen-strengthened

austenitic stainless steels at temperatures between 4 and 300 K [25].

3.3. Fatigue

The strain cycling fatigue behavior of 21Cr-6Ni9Mn has been measured at 295, 76 and 4 K by

Shepic and Schwartzberg [14] . The results at 4 K,

shown in fig. 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 nitrogen-strengthened grade does not result in a comparable improvement in fatigue life. The fatigue crack growth behavior of several nitro- gen-strengthened grades has been studied at cryogenic

temperatures [25,28]. The results for the high nitro- gen 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 4 K 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—6Ni—

9Mn. For the 21 Cr— 12Ni—5Mn alloy, the fatigue crack growth rates fall within the same scatter band

as the 300 series (A^< 0.08%) alloys at 4 K. For the Cr—Ni alloys, comparison of 304 and 304L with 304N

and 304LN shows that the nitrogen-strengthened

25
//./. McHenry, R.P. Reed / Structural alloys for superconducting magnets

alloys have growth rates at 4 K 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 4 K as the 316LN alloy.

3.4. Physical properties

In general, the physical properties of the nitrogenstrengthened stainless steels are similar to the 300-

series alloys, as summarized in table 3. In general, the nitrogen contributes to austenite stability and

thus reduces the magnetic permeability induced in some alloys by deformation at low temperatures.

Anomalous behavior has been observed in high manganese 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 sec- tion.

4.1. Welding

The austenitic stainless steels, with and without

nitrogen strengthening, are readily weldable by all of the common welding processes, providing the appropriate consumables and procedures are used. The service experience with welded assemblies of the

alloys has been satisfactory at cryogenic temperatures.

The American Welding Society, AWS, designation sys- tem, which is similar to the AISI system for wrought

products, is generally used to classify stainless steel

filled 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 4 K generally ex- ceeds the corresponding base metal strength, but toughness is usually significantly lower. Three pheno- mena 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

Table 5 Compositions of filler metals for welding austenitic stainless steels AWS type covered electrodes Composition (%)

Cr Ni C, max Other

E308 18 -21 9 -11 0.08 E308L 18 -21 9 -11 0.04 E310 25 -28 20 -22.5 0.20 0.75 Si, max

E16-8-2 14.5-16.5 7.5- 9.5 0.10 0.50 Si, max: E316 17 -20 11 -14 0.08 2-2.5 Mo

E316L 17 -20 11 -14 0.04 2-2.5 Mo

E330 14 -17 33 -37 0.25

Rod and bare electrodes ER308 19.5-22 9-11 0.08 ER308L 19.5-22 9-11 0.03 ER310 25 -28 20-22.5 0.08-0.15 ER316 18 -20 11-14 0.08 2-3 Mo

ER316L 18 -20 11-14 0.03 2-3 Mo

For covered electrodes: Mn = 2.5, max; Si = 0.9, max; P = 0.04, max; S = 0.03 max.

For rod and bare electrodes: Mn = 1.0— 2.5; Si = 0.25—0.60; P = 0.03; S = 0.03, max.

26
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

generally has a stronger influence on the weld metal

than the base metal.

Sensitization, the grain boundary precipitation of

Cr carbides, reduces weld metal toughness at low

temperatures [35]. Consequently, weld metals with

extra low carbon, such as 308L and 316L, are com- monly used for cryogenic applications. Sensitization

is most extreme if the weldment is given prolonged

exposure in the temperature range 500—800°C. However, the underlying weld beads in multipass weld- ments may receive sufficient exposure to these temperatures to cause significant toughness reductions

in stainless steels with greater than 0.03% carbon.

Subsequent annealing of the welds at temperatures

greater than 950°C dissolves the carbides and im- proves the toughness.

The ferrite content of as-deposited weld metals can be estimated by using the Schaeffler constitu- tion diagram. To avoid microfussuring the weld

metal chemistry is generally balanced to provide

4— 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 76 K for a wide range of stainless steel weld metals. Read,

McHenry and Steinmeyer [35] have evaluated the

Fig. 9. Fracture toughness of AWS 316L weld metal at 4 K

as a function of ferrite content [35].

effect of ferrite content on the fracture toughness

of shielded metal arc welds of AWS 316L; the results

at 76 and 4 K are shown in fig. 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—0.13% may occur depending on the amount of air which enters the arc shielding

medium. The greatest variations in nitrogen pickup

are associated with gas metal arc welding where the nitrogen 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 nitrogen pickup than titania 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— 1 100°C following severe forming operations

and stress relief at 420—480°C 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 76 K in types 308 and 310, but does not affect 308L. Stress relieving at 840°C followed by furnace

cooling causes about a 50% toughness loss in each

alloy (20-32 J as-welded and 8—20 J stress-relieved). Stress relieving at 950°C 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 1070°C 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 76 K 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 AISI types 304 and 316 are ACI

(Alloy Casting Institute) types CF8 and CF8M, respect27
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

ively. 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 the use of higher nickel content alloys such as ACI

grades CK20 (corresponding to AISI grade 310). A

16Cr—20Ni-10Mn-2.2Mo alloy has been used for liquid hydrogen [41] and liquid helium bubble

chambers [42,43] ; the alloy retains excellent strength (550 MPa yield strength) and ductility at temperatures to 2.5 K [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 conductivity. 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

21 Cr—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 coil 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 highly rated [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 large-grain-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 principal 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 low strength in weld- ments, high electrical and thermal conductivity, and

an unfavorable elastic modulus (too low) and thermal

expansion (too high) for many applications. For fusion magnet structures the preferred alumi- num 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

1 7 1 MPa (25 ksi) at 4 K. Alloy 2219 is generally used

in a heat treated condition (T85 1 , T87, etc.), and in the T851 temper has a typical yield strength of 482

MPa (70 ksi) at 4 K. Alloy 6061 is usually used in the heat treated condition (T6, T651 , etc.), and in the T6 temper has a typical yield strength of 380 MPa

(55 ksi) at 4 K. 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 4 K and may be used for special applications. Commercially pure aluminum (alloy 1 100)

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.

28
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Table 6 Compositions of aluminum alloys suitable for liquid helium service [46]

Type Composition (%)

Si Fe Cu Mn Mg Cr Zn Ti Others max max max

1100 1.0 Si + Fe, max 0.05-0.20 0.05, max — — 0.10 — 99.00 Al, minimum

2219 0.2 0.3 C/1 00 1On 00 0.20-0.40 0.02, max — 0.10 0.02-0.10 0.05-15 V;

0.10-0.25 Zr 3003 0.6 0.7 0.05-0.20 1.0 -1.5 - - 0.10 -

5083 0.4 0.4 0.1, max 0.40-1.0 4.0-4.

9

0.05-0.25 0.25 0.15, max

6061 0.4-0.

8

0.7 0.15-0.40 0.15, max 0.08-1.2 0.04-0.35 0.25 0.15, max

The ASTM specifications for aluminum mill pro- ducts 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 sec- tion the general trends in strength, toughness and

fatigue resistance are summarized, particularly the data for room remperature and 4 K. 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).

5.1. Tensile properties

The mechanical properties of alloys 5083, 6061

and 2219 at room temperature and 4 K 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 4 K and the ductility values are not significantly changed. The alloys

retain good notched-tensile properties at 4 K. Aluminum alloys having a wide range of strength

levels are suitable for liquid helium applications. The

yield strengths of commercially pure (CP) aluminum

(types 1 100-0 and 1099-H14) and the preferred alloys

are shown as a function of temperature in fig. 10 [42].

The yield strengths of each alloy gradually increase with decreasing temperature. Strength increases in the following order: annealed CP (1 100-0), cold

Table 7 ASTM specifications and availability of aluminum mill products [46]

Product form Specification Available alloys ASTM 5083 6061 2219

Sheet and plate B209 X X X Wire, rod, bar; rolled or B21

1

X X

cold finished Wire, rod, bar, shapes, tube B221 X X X

Structural shapes B308 X

Pipe B24

1

X

Forgings and forging stock B247 X X X

Structural pipe and tube; B249 X

extruded

29
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Table 8 Tensile properties of aluminum alloys at room temperature and 4 K [50]

Alloy and

temper

Plate thickness (mm)

Orientation Temp.

(K) TStensile strength

(MPa)

YS

yield strength

(MPa)

Elongation

in 4D

(%)

Reduction

of area (%)

NTS

notch

tensile strength (MPa)

NTS

TS NTS

YS

5083-0 25 L RT 322 141 19.5 26 372 1.16 2.65 4 557 178 32 33 429 0.77 2.42 5083-H321 25 L RT 335 235 15 23 421 1.26 1.80 4 591 279 29 33 508 0.86 1.82 6061-T651 25 L RT 309 291 16.5 50 477 1.54 1.64 4 483 379 25.5 42 619 1.28 1.63 T RT 309 278 15.2 42 467 1.51 1.68

4 485 363 23 33 601 1.24 1.66 2219-T851 25 L RT 466 371 11 27 547 1.12 1.48 4 659 484 15 26 703 1.06 1.48 T RT 457 353 10.2 22 531 1.16 1.50

4 659 481 13 21 665 1.01 1.38 2219-T87 25 L RT 464 387 11.8 28 567 1.22 1.46 4 674 511 15 23 690 1.03 1.37 A356-T61 19 casting RT 287 208 8.8 10 354 1.23 1.70

casting 4 455 331 7.1 9 495 1.09 1.50

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

TEMPERATURE, K

Fig. 10. Yield strengths of six aluminum alloys at temperatures between 4 and 300 K [50]

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 4 K 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—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 alu- minum alloys at 4 K. Tobler and Reed [53] measured

the fracture toughness of 5083-0 (4.3 cm thick plate)

30
H.I. McHenry, R.P. Reed /Structural alloys for superconducting magnets

Table 9

Physical properties of aluminum alloys

Property Units 5083 (annealed) Precipitation hardened 606 1 Precipitation hardened 2219

295 K 4 K Ref. 295 K 4 K Ref. 295 K 4 K Ref.

Specific gravity g/cm 3 2.66 2.70 2.83 Young’s modulus GN/m2 71.5 80.9 58 70.1 77.7 7 77.4 85.7 59

Shear modulus GN/m2 26.8 30.7 58 26.4 29.2 7 29.1 32.5 59

Poisson’s ratio 0.333 0.318 58 0.338 0.327 7 0.330 0.318 59 Thermal conductivity W/m • K 120 3.3 7 120-160 3-6 50 120-160 3-6 50 Thermal expansion K

-i xl0~6 23 14.1 7 120-160 14.1 7 120-160 14.1 7 Specific heat J/kg • K 900 0.28 7 120-160 0.28 7 0.28 0.28 7

Electrical resistivity ti fl cm 5.66 3.03 17 3.94 1.38 17 5.7 2.9 17

using /-integral procedures and found that toughness

increased with decreasing temperature. The A^ c (/) was 47 MPa\/m at 4 K in the transverse (TL) orientation. Shepic [54] measured the toughness of 2219-

T87 (3.8 cm thick plate) using the ASTM E399 plane

strain fracture toughness testing procedures. The K\ c for 2219-T87 at 4 K was 45.5 MPa\/m in the trans- verse (TL), 47.7 MPa\/iTi 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, 77 and 4 K. For a given strain

level, the fatigue lives of both alloys were greater at 4 K than at the higher temperatures for endurances

o

Fig. 1 1. Fatigue behavior of 2219-T851 aluminum alloy at temperatures of 300, 77 and 4 K [13].

greater than about 500 cycles. The results for 2219-

T851 are shown in fig. 1 1. The Manson-Hirschberg

[55] method of universal slopes gave reasonable pre- dictions 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] alloys have been

determined at 4 K. In each case the results indicate

that the growth rates at 4 K are slower than the rates

at room temperature.

5.4. Physical properties

The physical properties of the candidate aluminum

alloys are summarized in table 9. The thermal properties have been taken from the best -fit 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 resistivity [17],

With the exception of magnetic susceptibility, the physical properties of aluminum alloys are not

as good as those of stainless steels for fusion magnet

applications at 4 K. 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 increased 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

31
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

Table 10 Compositions of aluminum welding rods and electrodes for liquid helium service [46] ; composition, % (max unless range given)

Type Cu Mg Mn Cr Si Fe Zn Ti Be Other Suggested

for welding

1100 0.20 - 0 05 - Si + Fe = 1.0, max 99.0 -99.5 A1 1100, 3003 2319 5. 8-6.

0.02 0.20-0.40 - 0.20 0.30 0.10 0.10-0.20 0.0008 0.10- 0.25 Zr 2219 4043 0.03 0.05 0.05 - 4. 5-6.0 0.80 0.10 0.20 0.0008 0.05- 0.15 V 1100,6061 5183 0.10 4. 3-5.

0.50-1.0 0.05-0.25 0.40 0.40 0.25 0.15 0.0008 5083 5356 0.10 4. 5-5.5 0.05-0.20 0.05-0.20 Si + Fe = 0.5, max 0.10 0.06-0.20 0.0008 5083,6061

5556 0.10 4. 7-5.5 0.5 -1.0 0.05-0.20 Si+ Fe = 0.4, max 0.25 0.05-0.20 0.0008 5083,6061

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 received 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

Table 1

Tensile properties of aluminum alloy weldments at room temperature and 4 K (60)

Alloy and

temper

Plate Filler thichness alloy (mm)Post- Temp,

weld (K)

thermal

treat- ment

TS

tensile strength

(MPa)

JYS

joint

yield strength (MPa)

Elongation in 4 D

(%)

Reduction of

area (%)

Joint

effi- ciency (%)

Location of

failure

NTS

notch

tensile strength (MPa)

NTS

TS NTS

JYS

5083-0 25 5183 no RT 293 138 21.5 _ 100 WM 308 1.05 2.22

4 381 174 27 37 69 BM 371 0.98 2.14 5083-H321 25 5183 no RT 305 179 14 39 96 HAZ 376 1.23 2.10 4 455 246 9.0 14 80 WM 405 0.89 1.65 5083-H321 25 5356 no RT 286 167 13.5 47 90 WM 371 1.30 2.22

4 455 235 9.0 17 80 HAZ 398 0.87 1.69 5083-H321 25 5556 no RT 306 176 14 36 97 WM 370 1.21 2.10 4 474 238 13 17 83 WM 399 0.84 1.68 6061-T6 25 4043 no RT 214 144 6.0 19 69 HAZ 234 1.10 1.63 4 338 259 4.5 9 63 WM 275 0.81 1.06 yes RT 298 247 11 44 96 BM 396 1.31 1.57 4 452 309 15 16 84 WM 463 1.02 1.50 6061-T6 25 5356 no RT 225 156 8.0 31 73 WM 323 1.44 2,07 4 398 243 13.5 24 74 WM 367 0.92 1.50 yes RT 279 202 9.5 33 90 BM - - -

4 476 307 19 24 89 WM 419 0.88 1.37 2219-T851 32 2319 no RT 225 185 2.0 5 50 HAZ 280 1.24 1.52 4 411 277 2.5 10 62 HAZ 363 0.88 1.31 2219-T62 25 2319 yes RT 395 277 7.5 7 99 HAZ 439 1.11 1.58 4 496 355 3.5 5 80 HAZ 570 1.15 1.61

Note: Location of failures: WM - weld metal; BM - base metal; HAZ - heat affected zone.

32
H.I. McHenry, R.P. Reed /Structural alloys for superconducting magnets

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 4 K are summarized in table 1 1 . 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 alu- minum 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 1 1 also indicate that the 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 4 K; and the best ductility and notch-yield ratios in the as-welded

condition at 4 K 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: O-annealed, H-strain hardened

and T-thermally treated. For the strain hardened

tempers, the first digit indicates the hardening operations: HI -strain hardening only, H2-strain hardened

and partially annealed and H3-strain hardened and

stabilized. The second digit (e.g. HI 8 is strain-harden- ed 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 sys- tem.

6.3. Castings

Aluminum alloy castings can be produced by a variety of processes and generally are satisfactory for nonstructural 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 nonwelded

structural applications. Nelson, Kaufman and Wanderer

[60] evaluated the tensile properties at 4 K 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.

6.4. Machining, forming and forging

Machining'. Aluminum alloys, particularly the heat

treated grades, have excellent machinability. 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 4 K 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 370-5 10°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 conventional methods: dye-penetrant,

ultrasonics and radiography.

33
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

7. Alternative alloys

Alternative alloy systems do not provide candidate

structural materials for superconducting magnets in fusion energy systems. Consider the following sys- tems:

Nickel base alloys are ferromagnetic, expensive and

frequently difficult to fabricate. However, many of

these alloys have excellent combinations of strength

and toughness at 4 K and may be suitable for special

applications [61 ]

Titanium alloys are brittle at 4 K, expensive and

difficult to fabricate [62].. Cobalt base superalloys transmute to long half-life reaction products and are expensive and difficult to

fabricate. Carbon and alloy steels including the ferritic and

martensitic stainless steels are ferromagnetic and

brittle at 4 K. Even the cryogenic nickel steels are

brittle at 4 K [63] except for a specially processed

1 2Ni alloy [64].

Copper alloys are acceptable, but generally not competitive with the austenitic stainless steels or aluminum alloys.

8. Summary

This review suggests that there are several commercially available alloys that are suitable for construction of large superconducting magnet structures

for fusion energy systems. These alloys include anneal- ed austenitic stainless steels, nitrogen strengthened

austenitic stainless steels where greater strength is needed, an annealed Al-Mg alloy for welded alumi- num construction and precipitation-hardened alumi- num alloys for bolted assemblies. For each of these

alloys there are enough experimental data and fabrica- tion 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.

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 4 K,

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. Estab- lishment 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 tech- nology 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, a descrip- tion of the operating loads should include the anti- cipated load history for normal operation, loads due

to credible fault conditions and procedures for treat- ing combined loads. Tire 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 suffi- cient to preclude the use of materials in the magnet

systems which have properties inferior to those as- sumed in design. Inspection specifications and stan- dards are needed to provide a high degree of confi- dence that no defects are present in the raw materials,

or introduced during fabrication, which could cause

unsatisfactory performance of the magnet system.

34
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

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[ 1 ] R.P. Reed, A.F. Clark and E.C. van Reuth, in: Advances

in Cryogenic Engineering, vol. 22 (Plenum Press, New

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in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures, vol. 1, NBSIR 78-884

(National Bureau of Standards, Boulder, CO, 1978)

243.

[3] F.R. Fickett, M.B. Kasen, H.I. McHenry and R.P. Reed,

in: Advances in Cryogenic Engineering, vol. 24 (Plenum

Press, New York, 1979).

[4] F.R. Fickett, Properties of Non-superconducting Technical Solids at Low Temperatures, Proc. Fourth International Conference on Magnet Technology Brookhaven, New

York, 1972) 498.

[5] H.I. McHenry, in: Advances in Cryogenic Engineering,

vol. 22 (Plenum Press, New York, 1976) 9. [6] Handbook on Materials for Superconducting Machinery,

Metals and Ceramics Information Center, MCIC-HB-04

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[7] D.B. Mann, LNG Materials and Fluids (National Bureau

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[8] K.G. Brickner and J.D. Defilippi, in: Handbook of

Stainless Steels (McGraw-Hill, New York, 1977) 20.

[9] Low Temperature and Cryogenic Steels, Materials Manual (U.S. Steel Corp., Pittsburgh, PA, 1966).

[10] Materials for Cryogenic Service-Engineering Properties of Austenitic Stainless Steels (International Nickel Limited, London, 1974).

[11] ASME Boiler and Pressure Vessel Code, Section VIII, Pressure Vessels, (American Society of Mechanical

Engineers, New York, 1977).

[12] R.L. Tobler, in: Fracture 1977, vol. 3, Proc. Fourth

International Conference on Fracture (Waterloo, Canada,

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[13] A.J. Nachtigall, in: Properties of Materials for Liquefied

Natural Gas Tankage, ASTM STP 579 (American Society

for Testing and Materials, 1975) 378.

[14] J.A. Shepic and F.R. Schwartzberg, in: Materials Studies

for Magnetic Fusion Energy Applications at Low Temperatures-1, NBSIR 78-884 (National Bureau of Standards, Boulder, CO, 1978) 13. [15] R.L. Tobler and R.P. Reed, in: Advances in Cryogenic

Engineering, Vol. 24 (Plenum Press, New York, 1979).

[16] H.M. Ledbetter, W.F. Weston and E.R. Naimon, J. Appl.

Phys. 46 (9) (1975) 3855.

[17] A.F. Clark, G.E. Childs and G.H. Wallace, Cryogenics

10 (1970) 295.

[18] K.R. Efferson and W.J. Leonard, Magnetic Properties of Some Structural Materials Used in Cryogenic Applications; ORNL-4 150 (Oak Ridge National Laboratory,

Oak Ridge, TN, 1976) 126.

[19] Metals Handbook, vol. l,8thedn. (American Society

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[20] G.P. Sanderson and D.T. Llewellyn, J. Iron Steel Inst. Lond. 27 (1969) 1129.

[21] V.A. Randak, W. Wesseling, H.E. Bock, H. Steinmaurer and L. Faust, Stahl Eisen 91 (1971) 1255.

[22] R. Voyer and L. Weil, in: Advances in Cryogenic Engineering, vol. 1 1 (Plenum Press, New York, 1966) 447.

[23] D.C. Larbalestier, Private Communication, University of Wisconsin, Madison, WI (1977).

[24] C.E. Spaeder, Jr. and W.F. Domis, Project 44.012-062(2),

US Steel Corp., Pittsburgh, PA (1969).

[25] D.T. Read and R.P. Reed, in: Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-1,

NBSIR-78-884 (National Bureau of Standards, Boulder, CO, 1978) 91.

[26] D.B. A. Macmichael, M. Liddle and J.S.H. Ross, Superconducting AC Generators Trial Rotor Forging Investiga- tion (International Research and Development Co. Ltd., Newcastle upon Tyne, England, 1978).

[27] D.T. Read and R.P. Reed, in: Materials Studies for Magnetic Fusion Energy Low Temperature Applications-II

(National Bureau of Standards, Boulder, CO, 1979).

[28] R.L. Tobler and R.P. Reed, in: Elastic Plastic Fracture, ASTM STP-000, (American Society for Testing Materials, Philadelphia, PA, forthcoming).

[29] P.F. Scardigno, M.S. Thesis, Naval Postgraduate School, Monterey, CA, NTIS AD/A-004555 (1974).

[30] P.R. Landon, Unpublished Data, Lawrence Livermore

Laboratories, Livermore, Ca (1975).

[31] R.D. Masteller, NASACR-82638, Martin Marietta Corp., Denver, CO (1970).

[32] C.O. Malin, NASA SP-5921 (01), Technology Utiliza- tion Office, NASA, Washington, D.C. (1970).

[33] H.M. Ledbetter, Mater. Sci. Eng. 29 (1977) 255.

[34] H.M. Ledbetter and E.W. Codings, in: Metal Physics of

Stainless Steels (TMS AIME, New York, forthcoming).

[35] D.T. Read, H.I. McHenry and P.A. Steinmeyer, in: Materials Studies for Magnetic Fusion Energy Low

Temperature Applications, vol. II (National Bureau of

Standards, Boulder, CO, 1979).

[36] E.R. Szumackowski and H.F. Reid, Weld. J. 57 (1978)

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(1950) 493-s. [39] J.H. McConnell and R.R. Brady, Chem. Eng. 67 (1960)

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vol. 20 (Plenum Press, New York, 1965) 26. [42] D.F. Shaw, Cryogenics 4 (1964) 193. [43] A.B. Miller, Cryogenics 5 (1965) 320.

[44] Armco Nitronic 40 Stainless Steel Sheet and Strip (Armco Steel Corp., Baltimore, MD, 1976).

[45] J.E. Jenson, ed., Forging Industry Handbook (Forging

Industry Association, Cleveland, OH, 1966).

[46] Aluminum Standards and Data, Aluminum Association,

Inc. Washington, D.C. (1976).

[47] H.I. McHenry, Met. Progr. 93 (3) (1968) 62. [48] J.G. Kaufman and E.T. Wanderer, in: Aluminum, vol. II: Design and Applications (American Society for Metals, Metals Park, OH, 1967) 297.

35
H.I. McHenry, R.P. Reed / Structural alloys for superconducting magnets

[49] Alcoa Aluminum — The Cryogenic Metal, Alcoa, Pitts- burgh, PA (1971).

[50] J.G. Kaufman, K.O. Bogardus and E.T. Wanderer, in: Advances in Cryogenic Engineering, vol. 13 (Plenum

Press, New York, 1968) 294.

[51] J.G. Kaufman and E.W. Johnson, in: Advances in Cryogenic Engineering, vol. 8 (Plenum Press, New York,

1963) 678.

[52] M.P. Hanson, G.W. Stickley and H.T. Richards, in: Low

Temperature Properties of High-Strength Aircraft and

Missile Materials, ASTM STP 287 (American Society

for Testing and Materials, Philadelphia, PA, 1961) p. 3. [53] R.L. Tobler and R.P. Reed, J. Eng. Mat. Technol. 99

(1977) 306.

[54] J.A. Sphepic, in: Materials Studies for Magnetic Fusion Energy Low Temperature Applications, vol. II (National Bureau of Standards, Boulder, CO, 1979). [55] S.S. Manson, Exp. Mech. 5 (1965) 193. [56] H.I. McHenry, S.E. Naranjo, D.T. Read and R.P. Reed,

in: Advances in Cryogenic Engineering, vol. 24 (Plenum

Press, New York, 1979).

[57] G.E. Childs, L.J. Ericks and R.L. Powell, Thermal Conductivity of Solids at Room Temperature and Below,

NBS Monograph 131 (National Bureau of Standard,

Boulder, CO, 1973).

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Sci. 10 (1975) 1309.

[59] D.T. Read and H.M. Ledbetter, J. Eng. Mater. Technol. 99 (1977) 181. [60] F.G. Nelson, J.G. Kaufman and E.T. Wanderer, in: Advances in Cryogenic Engineering, vol. 14 (Plenum Press, New York, 1969) 71. [61] J.M. Wells, R. Kossowsky, W.A. Logsdon and M.R. Daniel,

in: Materials Research for Superconducting Machinery

VI, NTIS No. ADA 036919 (National Bureau of

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36
STRENGTH AND TOUGHNESS RELATIONSHIP FOR INTERSTITIALLY STRENGTHENED

AISI 304 STAINLESS STEELS AT 4 K

R. L. Tobler, D. T. Read, and R. P. Reed

National Bureau of Standards

Boulder, Colorado

ABSTRACT

A study was conducted to determine the effects of carbon and nitrogen

on the 4-K fracture properties of Fe-18Cr-10Ni austenitic stainless

steels having a base composition corresponding to AISI 304. J-integral

fracture toughness tests using 24. 5-mm-thick compact specimens (TL

orientation) were performed in a liquid helium environment on nine steel

heats having carbon plus nitrogen (C+N) contents between 0.067 and

0.325 wt.%. The fracture toughness decreased with increasing C+N content.

The Kj estimates obtained at 4 K ranged from 337 to 123 MPa*m^ 2

,

exhibiting an inverse dependence on tensile yield stress. A computeraided J-integral test facility was implemented to conduct this study.

This new facility improves measurement accuracy, conserves material

specimens and testing time, and systematizes test procedures.

37

INTRODUCTION

Acceptance of the parameter J^

c as a measure of the fracture resistance of metals has become widespread in recent years. A large volume

of J-integral fracture toughnesss data on a variety of candidate materials

for low temperature applications has been accumulated. The J-integral

test has the advantage of being applicable to smaller specimens (up to

100 times smaller) and tougher materials than the plane strain fracture

test (Kj

c test). A disadvantage of the J-integral technique has been

its experimental difficulty. Previous test methods have required approximately

five specimens and lengthy data reduction procedures [1]. Recently

developed single-specimen J-integral test techniques [2] eliminated the

multiple-specimen requirement, but the data reduction procedure was

still lengthy. In the computer-aided J-integral test, most of the datareduction procedure is automated; the apparatus produces a plot of the

J-integral value versus crack extension from which Jj

c can be quickly

and simply extracted.

The AISI 300 series stainless steels are austenitic Fe-Cr-Ni

alloys, offering relatively low strength but excellent cryogenic ductility and toughness [3]. By virtue of favorable mechanical properties,

service history, and availability, AISI 304 is perhaps the most widely

used cryogenic alloy in the world. But ni trogen-strengthened grades

such as AISI 304 N or AISI 304 LN have attracted attention recently as

possible substitutes for AISI 304 in applications demanding higher

strength. Nitrogen is a relatively inexpensive and effective strengthener

Nitrogen also stabilizes the austenitic structure and reduces the probability

of martensitic transformation when stress is applied at cryogenic temperatures

39
Carbon has similar effects, but carbon contents in austenitic stainless

steels must be held to low levels to prevent sensitization (chromium

carbide precipitation) during thermal excursions.

Nitrogen-strengthened austenitic stainless steels are gaining

popularity, but some applications are hindered by limited availability,

limited service experience, lower fracture toughness, and lack of

cryogenic design data. Several mechanical property studies were conducted on these steels [4-18], but most papers are concerned only with

room temperature behavior. In the present study, fracture toughness

measurements for nine heats of an Fe-1 8Cr-l ONi stainless steel are

reported at 4 K. The carbon and nitrogen (C+N) contents of the nine

heats were varied to enable a systematic study of interstitial concentration.

The results can be used to predict the strength and toughness of commercial

steels of given compositions or to tailor the properties for a particular

application by appropriate specification of the steel composition.

MATERIALS

Nine stainless steel plates were procured from the research laboratories of a major steel manufacturer. These plates had a nominal

base composition falling within the limits set by ASTM specification

A 240 for AISI 304 stainless steel: Fe-1 8Cr-l ONi-1 . 5Mn-0 . 02P-0 . 02S-0 . 55Si

0.2Mo-0.2Cu. The carbon and nitrogen levels varied, however, with

carbon at 0.03, 0.06, or 0.09 wt .% and nitrogen at 0.04, 0.12, or 0.24

wt.%. The mill chemical analyses are listed in Table I. The nine

plates were produced from three 1 36- kg (300-lb) vacuum-inductionmelted heats, split with respect to carbon level, and teemed into 76- x

200- x 360-cm hot-topped cast-iron ingot molds. The ingots were then

reheated and soaked at 1561 K (2350°F), hot-rolled to 25 . 4-mm-thi ck plates,

40
and air cooled. The plates were finally annealed at 1332 ± 7 K (1937° ±

13°F) for 1 h and water quenched. The hardness and grain size measurements

are 1 isted in Tabl e I

.

METHODS

Since elastic-plastic fracture was anticipated in the test materials,

the Jj

c parameter was measured. The measurement problem in Jj

c testing is to obtain the required plot of J-integral, J, versus crack

extension, Aa , from the directly measured quantities, which are the load, P,

on the specimen and the relative displacement, 6 , of two measurement points

located on the specimen load line on opposite sides of the notch (Fig. 1).

The J value is proportional to the area under the (P,<5) curve (Fig. 2) and

can be easily determined for any point along the curve. The shaded

area in Fig. 2 is proportional to J at the point labeled "a". The principle

used to measure Aa is that the crack length, a, is related to the slope

of the ( P , 6 ) curve of the specimen for elastic conditions through the

known specimen compliance function. The Aa measurement procedure is to partially

(about 10%) unload the specimen while recording load displacemnt changes

(Fig. 2 and 3). Subsequent unloadings at increasing crack extension

values result in increased compliance. The crack extension at a given unloading is the difference between the crack length measured in that unloading

minus the initial crack length measured by an unloading carried out before

crack extension.

The cryostat system that was used for computer-aided J-integral testing

was described previously by Fowlkes and Tobler [19]. The specimen

loading is controlled in conventional fashion by the operator. The computer

41
is used only in data acquisition and reduction. An analog ramp generator

controlled by the operator applies a smoothly varying signal to the

test-machine control circuitry to allow unloading-reloading cycles as

necessary. The additional equipment needed for computer-aided testing

includes a multichannel analog-digital conversion system, the computer

itself, a CRT terminal, a dual floppy disk data-storage unit, and a

digital plotter (Fig. 4).

Two signals are used to direct the operation of the data acquisition

system. One is derived from the output of the ramp generator by which

the operator initiates the unloading-reloading cycle. When the ramp

generator is activated, a signal transmitted to the processor commands

it to begin to acquire and store load- and displacement-change data

during the unloading-reloading cycle. Data are stored -until the load

again reaches the value it had before the unloading. The test continues

through as many unloading cycles as desired until the operator activates

the quit signal, which directs the processor to await the operator's

command, entered at the CRT terminal, to stop the data acquisition

procedure and terminate the test.

During the incremental unloading required to measure Aa, displacement

changes of the order of 0.05 mm and load changes of the order of 5 kN

must be correlated accurately. These changes correspond to a change in

the overall displacement signal of about 70 mV and a change in the load

signal of about 1 V. Because the analog-to-digi tal converter used has

only 16-bit resolution, the displacement signal cannot be converted

directly to digital form with precision sufficient for accurate measurement of this displacement change. The measurement of the small displacement changes that occur during the unloading-reloading cycles is

42
facilitated by regarding the overall displacement to consist of small

changing displacements superimposed on a large reference displacement

value. The measurement is accomplished by canceling out the reference

part and then converting the small displacement changes to digital form

with high resolution. To achieve this cancellation, the overall displacement is converted to digital form at the beginning of each unloadingreloading cycle. This initial value is treated as the reference part.

To cancel out this reference part of the displacement signal, a digitalto-analog converter with 14-bit resolution is used to apply a voltage

equivalent to the reference part of the displacement signal to one

terminal of a differential high-sensitivity anal og-to-digi tal conversion

channel. The overall displacement signal is applied to the other input

terminal of this channel. This results in the anl aog-to-digi tal conversion

of a voltage equal to the difference between the instantaneous value of

the displacement signal and its value at the beginning of the unloading.

This difference signal is the required change-of-di spl acement signal and

is converted to digital form with 16-bit resolution of its 80-mV range.

The least significant bit represents about 2.5 yV, which is quite sufficient.

The load-change signal is produced in a similar manner by using another

digital -to-analog converted signal and converted to digital form with

16-bit resolution at the 1.25-V level. Here the least significant bit

represents about 40 yV, which is again quite sufficient.

Approximatel y 80 load-change and displacement-change data points

are obtained for each unloading-reloading cycle over a 30-s period.

These data are plotted during the unloading-reloading cycle with a

43
digital plotter, forming a plot like Fig. 3. From this plot, the

presence of excessive noise or other improper behavior in the signals

can be easily observed. The load-change and displacement-change data

are correlated using a least-squares fit. From the resulting slope, the

instantaneous crack length is calculated using the known expression for

the compliance of the specimen.

Each (J, Aa) datum is plotted by the digital plotter as soon as it

is available, forming a plot similar to Fig. 1. This plot allows convenient monitoring of the progress and quality of the test in real time.

The values are also displayed on a CRT terminal for transcription. This

step will be replaced by an automatic print-out when facilities become

available. The 1 east-squares-fi t and data-output procedure takes about

12 s.

The record of the completed test consists of the (J, Aa), (P, 6),

and (aP, A6 ) plots, the transcribed (J, Aa) data, and three files stored

on a floppy disk during the test—one containing the (J, Aa) values and

a few other key data, one containing all the (P, 6 ) data, and one containing

all the (aP, a 6 ) data.

The computer-aided J-integral test facility at NBS is similar to

that described by Joyce and Gudas [20]. The differences are in the

specific types of equipment used and in the real-time outputs. Realtime is used here to mean during the course of the test, as opposed to

after the test is completed.

The procedure for extracting J and Aa from the directly measured

load and displacement on a deeply notched, compact tensile specimen is

as follows: The electrical signals from the load cell and displacement

44
gage are amplified by the electronics supplied with the closed-loop

servo-hydraul ic apparatus used for the test. A plot of load versus displacement similar to Fig. 2 is produced by an x-y plotter throughout the

course of the test. Either signal may be displayed on a digital readout.

The plot and readout are used by the operator to monitor the progress of

the test.

The amplified load and displacement signals are also introduced

into an analog-to-digital conversion system to permit calculation of J

and Aa. At this point, the overall levels of the signals range between

0 and 10 V. These signals are converted to digital form with a resolution

of 16 bits; the least significant bit represents about 0.3 mV. The

overall load and displacement values are used for calculating the

instantaneous J value, which is proportional to the area under the loaddisplacement curve. Stored overal 1 -displ acement points are separated by

3

2.5 x 10 . Smaller observed overal 1 -displ acement increments are not

stored. Computations of the instantaneous J values are made when the

overall displacement has increased sufficiently from the previously

stored value.

The geometry of the compact specimens tested in this study is shown

in Fig. 5. Specimen thickness, B, was 24.5 mm, and width, W, was 50.8

mm. The specimens were machined in the TL orientation as defined by

ASTM Method E-399-74 [21]. The specimens were fatigue cracked at 4 K to

a crack length between 28.8 and 36 mm - 0.57 to 0.71). The specimens

were then J-tested in displacement control, and partial unloadings were

performed periodical ly. The J value at each unloading point was calculated

using the expression:

, _ aA

J Bb

45
where A is the Merkl e-Corten factor dependent on crack length [22],

A is the area under the load-versus-def 1 ection curve at the unloading

point, b is the specimen ligament (b = W - a), and B is the specimen

thickness. J-resistance curves consisting of many sets of J and Aa

points were generated, as shown by the example in Fig. 6.

The J T value for each steel was taken at the intersection of the

Ic

J-resistance curve with the blunting line. The blunting line is given

by J = 2a^A

a where is the flow stress, which is the average of the

yield and ultimate tensile stresses. As shown in Fig. 6, there is an

artificial offset of the data from the theoretical blunting line, so a

blunting line parallel to the theoretical blunting line but passing

through the data was used. Scatter in the J-resistance curves caused

uncertainties of about ±12% in the values of J T . I c The validity of Jj

c results was determined by the size criterion,

which requires that [1]:

d u 2 5 J T B, b, a >_ Ic a

f

In all low temperature fracture tests, partial transformation

the paramagnetic austenite to ferromagnetic a' (bcc) martensite was

observed (e.g., Ref. 23). The extent of a' transformation was estimated

by using a bar-magnet device [24]. Small amounts of e (hep) martensite

may also have formed, but this was not measured.

RESULTS AND DISCUSSION

The computer-aided J-integral methods have now been in use for nine

months and have proved quite sati sfactory . In comparison with the

46
multiple specimen and non-computer-aided single-specimen techniques previously used, the new methodology has eliminated about 4 h of data

reduction per test and has provided a real-time readout of the test

progress and quality. The cost of the complete data-acquisition system,

including analog-to-digital converters, processor, digital plotter, and

CRT terminal is estimated at about $25,000.

J T results were obtained for seven of the austenitic stainless

Ic

steels tested, but the data for the two lower strength heats (with C+N

0.067 and 0.097 wt.%) were invalid. The maximum J values in these tests

exceeded the allowable limits for 24. 5-mm-thick specimens. The resistance

curves deviated from the blunting line at exceptionally low J values,

and the resistance curves obtained could not be analyzed according to

conventional techniques because there was no obvious knee at the blunting

line-resistance curve intersection. The J T values for the seven steels

Ic

2 at 4 K range from 63 to 499 kJ/m , whereas the yield stresses for the

same materials at 4 K range from 329 to 1286 MPa (Table II). Higher

toughness in these seven steels is associated with lower interstitial

C+N contents, and hence lower strength. All of the steels tested

exhibited a ductile dimpled fracture mode at 4 K, as shown by the scanning

electron fractographs of Figs. 7 and 8.

Estimates of the plane strain fracture toughness parameter, K

j

c

,

which would be observed under linear-elastic conditions, were obtained

using the relationship:

where E is Young's modulus, v is Poisson's ratio, and Kj

c (J) denotes an

estimate of Kj

c from Jj

c

. The values of E and v for the stainless

47
steels at 4 K were estimated to be 206.8 GPa and 0.3 [25], respectively.

The calculated Kj estimates are plotted versus C + N content in Fig. 9

and versus yield strength in Fig. 10. As shown, the Kj

c (J) values

decrease linearly with interstitial concentration and are inversely

related to yield stress. Therefore nitrogen, which is a particularly

effective strengthener , also has a strong influence on fracture toughness

at cryogenic temperatures . An expression governing the relationship

between fracture toughness and yield stress was derived from Fig. 10:

K Ic

(J) = 500 - 0.3 ay

1/2

The data scatter of Kj (J) from this expression is ±20 MPa-m

The present results compare well with existing data, some of which

are also shown for comparison in Fig. 10. The AISI 304 N datum of Read

and Reed [5] falls slightly below the present data, reflecting differences

owing to metallurgy or measurement techniques. The Fe-Cr-Ni -Mn-N

steels tend to have higher nitrogen contents, higher 4-K strength, and

lower toughness compared with the Fe-Cr-Ni-N (AISI 304) steels [4].

Yet, the Fe-Cr-Ni -Mn-N steels in Fig. 10 appear to possess higher

toughness than would be expected from the extrapolated trend for FeCr-Ni-N steels. It is uncertain whether this is due to metallurgical

effects or measurement bias, since the Fe-Cr-Ni-Mn-N data were obtained

by other methods including the mul ti specimen technique, which have

greater Jj

c uncertainty.

The inverse relationship of a and K

T is well known for linear-elastic

y

fracture toughness data. Now there is appreciable evidence that the

inverse relationship holds in the elastic-plastic region as well.

48
Previous 4-K data for a variety of austenitic stainless steels and

uncontrolled metallurgical conditions [5] indicated a linear trend of

lesser slope than that shown in Fig. 10. It is advantageous that the

present data derive from a controlled alloy series in which strength was

varied by interstitial content while other compositional and processing

variables were constant.

The approximate amount of a' (bcc) martensite formed during fatigue

and fracture at 4 K is plotted as a function of interstitial content in

Fig. 11. These results provide only an approximate indication of the

extent of phase transformation, because the magnetometer measurements

were taken directly from the compact specimen fracture surfaces, which

are rough and unmachined. It is clear from Fig. 11 that some martensite

forms in all of the steels during testing.

The quantity of martensite that forms would be expected to increase

with the plastic-zone size. Referring to Fig. 11, more a' is detected

in steel specimens of lower strength, higher toughness, and presumably,

larger plastic-zone size. It is not clear at this time whether the

formation of martensite plays a significant role in the fracture process

of these metastable steels.

Stress-strain curves of various heats at 295, 195, 76 and 4 K are

presented in Figures 12-17. At 4 K, all heats exhibited discontinuous

yielding. Both the magnitude and the strain of initiation of the yield

drops increased with increasing austenite stability and strength. The

inflections of the curves at low temperatures result from the martensitic

transformation. Notice in Figure 12 that the effect of C + N at 4 K is

to increase the early deformation of austenite and that the ultimate

tensile strength remains relatively unchanged.

49
SUMMARY AND CONCLUSIONS

1. New computer-assisted J-integral test methods and facilities have

been developed and applied to cryogenic testing.

2. Fracture toughness measurements at 4 K are reported for Fe- 1 8Cr- 1 ON

i

austenitic stainless steels having C+N contents ranging from 0.067

to 0.325 wt . %

.

3. The fracture toughness at 4 K decreased with increasing C+N content,

and the Kj

c values estimated from Jj

c (ranging from 123 to 336

1/2 MPa*m '

) were inversely related to yield stress.

ACKNOWLEDGMENTS

This research was supported by the Naval Ship Research and Development

Center (J. Gudas, coordinator) and by the Office of Fusion Energy, DoE

(D. Beard, coordinator) . T. Whipple, A. Castillo, D. Burkhalter, and D.

Beekman assisted in performing the fracture tests.

50
REFERENCES

1. J. D. Landes and J. A. Begley, "Test Results from J-integral

Studies: An Attempt to Establish a Jj

c Testing Procedure," in

Fracture Analysis , ASTM STP 560, American Society for Testing and

Materials, Philadelphia, PA (1974), pp. 170-186.

2. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt,

"Single Specimen Tests for Jj

c Determination," in Mechanics of

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

Materials, Philadelphia, PA (1976), pp. 27-42.

3. R. P. Reed, F. R. Fickett, M. B. Kasen, and H. I. McHenry, "Magnetic

Fusion Energy Low Temperature Materials Program: A Survey,"

i n Materials Studies for Magneti c Fusion Energy Appl ications at Low

Temperatures-I , NBSIR 78-884, F. R. Fickett and R. P. Reed, Eds.,

National Bureau of Standards, Boulder, CO (1978), pp. 245-335.

4. D. T. Read and R. P. Reed, "Toughness, Fatigue Crack Growth, and

Tensile Properties of Three Nitrogen-Strengthened Stainless Steels

Cryogenic Temperatures," in Materials Studi es for Magneti c Fusion

Energy Applications at Low Temperatures-I , NBSIR 78-884, F. R.

Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder,

CO (1 978), pp. 91-154.

5. D. T. Read and R. P. Reed, "Fracture and Strength Properties of

Selected Austenitic Stainless Steels at Cryogenic Temperatures ,

i n : Material s Studi es for Magneti c Fusion Energy Appl ications at

Low Temperatures-I

I

, NBSIR 79-1609, F. R. Fickett and R. P. Reed,

Eds., National Bureau of Standards, Boulder, CO (1979), pp. 81-122.

51
6. D. T. Read, R. P. Reed, and R. E. Schramm, "Low Temperature Deformation

of Fe-18Cr-8Ni Steels," in: Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures-I

I

, NBSIR 79-1609, F. R.

Fickett and R. P. Reed, Eds., National Bureau of Standards, Boulder,

CO (June 1 979), pp. 151-172.

7. D. T. Read and R. P. Reed, "Heating and Strain Rate Effects in AISI

304 L Stainless Steel at 4 K," in: Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures-I

I

, NBSIR 79-1609,

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

Boulder, CO (1979), pp. 125-147.

8.. R. L. Tobler and R. P. Reed, "Tensile and Fracture Behavior of

Ni trogen-Strengthened, Chromium-Nickel -Manganese Stainless Steel at

Cryogenic Temperatures , " in Elastic-Plastic Fracture , ASTM STP 668,

American Society for Testing and Materials, Philadelphia, PA (1979),

pp. 537-552.

9. P. P. Dessau, "LN

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

Report No. TID/SNA-2083, Aerojet Nuclear Systems Company, Sacramento, CA

(1971).

10. D. C. Larbalestier and D. Evans, "High Strength Austenitic Stainless Steels for Cryogenic Use," in Proceedings of Sixth International

Cryogenic Engineering Conference , IPC Science and Technology Press

Ltd., Guildford, Surrey, England (1976), pp. 345-347.

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

Properties of Chromium-Nickel Stainless Steels," J . Mater

.

, JMLSA,

5 (1 970), pp. 413-430.

12. R. Voyer and L. Weil, "Tensile and Creep Properties of a High

Nitrogen Content 18/10 (AISI 304L) Stainless Steel at Cryogenic

Temperatures," in Adv. Cryo . Eng. 1

1

, K. D. Timmerhaus, Ed., Plenum,

New York (1965), pp. 447-452.

52
13. L.-A. Norstrom, "The Influence of Nitrogen and Grain Size on Yield

Strength in Type AISI 316L Austenitic Stainless Steel," Met. Sci .

,

21 (1977), pp. 208-212.

14. M. K. Booker and V. K. Sikka, "Effects of Composition Variables on

the Tensile Properties of Type 304 Stainless Steel," ORNL-5353,

Oak Ridge National Laboratory, Oak Ridge, TN (1977), pp. 1-38.

15. P. Soo and W. H. Horton, "The Effect of Carbon and Nitrogen on the

Short-Term Tensile Behavior of Types 304 and 316 Stainless Steels,"

Ward NA 3045-2, Westinghouse Electric Corporation, Madison, PA

(1973) pp. 1-48.

16. K. J. Irvine, D. T. Lewellyn, and F. B. Pickering, "High Strength

Austenitic Stainless Steels," J . Iron Steel Inst . , 1 99 (1961),

pp. 153-175.

17. K. J. Irvine, T. Gladman, and F. B. Pickering, "The Strength of

Austenitic Stainless Steels," J . Iron Steel Inst. , 207 (1969), pp.

1017-1028.

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

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

Report No. 44012-062-2, United States Steel Corporation, Pittsburgh,

PA (1969).

19. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a

Ti-6A1-4V Alloy at Liquid Helium Temperature, " Eng . Fract. Mech. ,

8 (1976), pp. 487-500.

20. J. A. Joyce and J. P. Gudas, "Computer Interactive Jj

c Testing of

Navy Alloys," in Elastic-Plastic Fracture , ASTM STP 668, American

Society for Testing and Materials, Philadelphia, PA (1979),

pp. 451-468.

53
21. "Standard Method of Test for Plane Strain Fracture Toughness of

Metallic Materials (Designation E399-74)," 1 974 Annual Book of

ASTM Standards , Part 10, American Society for Testing and Materials,

Philadelphia, PA (1 974), pp. 432-451.

22. 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 , 6 (1974), pp. 1-7.

23. R. P. Reed and C. J. Guntner, "Stress-Induced Martensitic Transformations in 18Cr-8Ni Steel," Trans. Metal! . Soc. AIME 230 (1 964),

pp. 1 71 3-1 720.

24. R. P. Reed and R. P. Mikesell, "The Stability of Austenitic Stainless Steels at Low Temperatures as Determined by Magnetic Measurements," in Adv. Cryo. Eng. , 4_, K. D. Timmerhaus, Ed., Plenum, New

York (1 960), pp. 84-1 00.

25. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, "ElasticConstant Variability in Stainless Steel 304," J . Appl . Phys

.

,

51 (1980), pp. 305-309.

54
LIST OF TABLES

Table No.

I. Stainless steel compositions (wt.%), hardnesses, and grain sizes.

II. Yield strength and J-integral toughness results for Fe- 1 8Cr- 1 ON

i

steels at 4 K.

55
lDCOOlOLDCOCOCMO cor-cnocor^cncocri

o cr> c\j lo co .— r~^ cn c\j OOr— r— r— C\J C\J C\J CO ooooooooo

= = OsJ = =

COCOO'NN^ONl/1 (NjincoroyDOininco ooooooooo

ooooooooo

• = = C\J = =

o= = o= = o= =

o = = o = =

o = = cr> = = cr* = =

oo= = co= = o= =

.— ojfO«tiniorscoo>

56
Table II. Yield strength and J-integral toughness results Fe- CO 0 -5 1 o 12 steels at 4 K.

Specimen C+N Content

(wt.%)

Yield Strength, o

(MPa)

a

Flow Strength,

(MPa)

25 J

(mm)

Ic

J Ic

( kJ/m^)

b

(MPa-m 1/2

) c

1 0.067 329 892 NA NA NA

2 0.097 445 996 NA NA NA

3 0.128 530 1029 12.2 499 337

4 0.157 745 1158 6.7 312 266

5 0.187 876 1288 4.5 230 230

6-A 0.214 896 1264 4.3 218 222

6-B 0.214 896 1264 5.3 272 249

7 0.270 1186 1411 1.2 67 124

8-A 0.297 1178 1456 1.5 88 141

8-B 0.297 1178 1456 1.3 78 133

9 0.325 1286 1496 1.1 63 124

2 ^Calculated from the expression K Ic

= Jj -E/(l -v ) , assuming E = 206.8 GPa and

V = 0.3 at T = 4 K.

a To convert from

I

j^To convert from

C To convert from

MPa to ksi, divide by 6.895.

2 2 kJ/m to in-lb/in , multiply by 5.709.

MPa'/m to ksi /in", divide by 1.099.

57
LIST OF FIGURES

Fi gure 1 .

Figure 2.

Figure 3.

Figure 4.

Figure 5.

Figure 6.

Figure 7.

Figure 8.

Figure 9.

Figure 10.

Fi gure 1 1

.

Figure 12.

Typical J-i ntegral -versus-crack extension curve.

Typical 1 oad-versus-specimen displacement curve, indicating J-i ntegral area for specimen at point a.

Calibration curve of load change versus displacement

change.

Flow chart of recording and analyses computer system for

fracture testing.

Specimen geometries used for fracture mechanics tests.

Typical J-resistance curve for an austenitic stainless

steel tested at 4 K.

Scanning electron microscopy photograph of Fe-1 8Cr-l ON i

;

C+N = 0.325 wt.%. The 4-K fracture surface is shown at

200 X and 1000 X.

Scanning electron microscopy photograph of Fe-1 8Cr-l ON i

;

C+N = 0.067 wt.%. The 4-K fracture surface is shown at

200 X and 1000 X.

Dependence of Kj

c on C+N content at 4 K.

Inverse relationship between fracture toughness and yield

stress for austenitic stainless steels at 4 K.

Relative amounts of a' (bcc) martensite formed during

tests of compact specimens at 4 K. (Arbitrary Units)

Engineering stress-stra i n curves for Fe-1 8Cr- 1 ON i steel

heat no. 1, representati ve of low interstitial C+N contents .

58
Figure 13.

Figure 14.

Figure 15.

Figure 16.

Figure 17.

Engineering stress-strai n curves for Fe- 1 8Cr- 1 ON i steel

heat no. 4, representative of medium interstitial C+N

contents

.

Engineering stress-strain curves for Fe- 1 8Cr-l ONi steel

heat no. 8, representative of high interstitial C+N

contents

.

True stress-strain curves for Fe-19Cr-10Ni steel heat

no. 1 (serrations at 4K are not shown).

True stress strain curves for Fe- 1 8Cr- 1 ONi steel heat

no. 4 (serrations at 4K are not shown).

True stress-strain curves for Fe-1 8Cr-l ONi steel heat

no. 8 (serrations at 4K are not shown).

59
009

(iu/N>o nvao3iNi-r

CD>

S-

=3U

CD

60

CRACK

EXTENSION,

Aa

(mm)
(n>d d ‘avcn

E

E

co

O

<

-j

CL

co

Q

"O

aj>

s_

rsu

c

cdE

CD

CD

raQ.

00

c

CD _Eu

CDCL

CO

CO

13CO

S_

CD>

I -a

rao

CD

CDQ.

>1

OJ

CD

S61

eating

J-integral

area

for

specimen

at

point
62

DISPLACEMENT

CHANGE,

AS

(mm)
CD

k- \

k<5

Digita Plotte

\s

0

)

<DO

O

CRT ermine

kQ

iCO

k_

o

=Q

a E

I

®

O cf) O >

CO cC/3

CO

•4—

»

’ 5) Q

D)

O

CO

c

o

CO

L»CD>

c

o

c o

<

"c3

E co 0 O +-* •«=» CO c

> o

CO - ^ o

CO ®

© UJ

63
COMPACT SPECIMEN

Thickness = 24.5 mm

Thickness = 2.54 mm

Figure 5. Specimen geometries used for fracture mechanics tests.

G4
1200

J-Resistance Curve

Fe-l8Cr-IONi Stainless Steel

(C+N = 0.128 weight percent)

Compact Specimen

24.5 mm thick

a/W» 0.630

TL Orientation

Blunting Line

J = 2Aao-f

0 12 3

Crack Extension, Aa (mm)

Figure 6. Typical J-resistance curve for an austenitic stainless

steel tested at 4 K.

65
Figure 7. Scanning electron microscopy photograph of Fe-1 8Cr-l OMi

;

C+N = 0.325 wt.%. The 4-K fracture surface is shown at

200 X and 1000 X.

n,c J -J
Figure 8. Scanning electron microscopy photograph of Fe-1 8Cr-l ON i

C+N = 0.067 wt.%. The 4-K fracture surface is shown at

200 X and 1000 X.

67
400

OJ

o

0=

300

—>"ofr—

a

in

in

0}

c. JZ

CP

3

o

I—

<D

k-3

o

o

200

100

0.10 0.20 0.30

Carbon + Nitrogen (weight percent)

Figure 9. Dependence of Kj on C+N content at 4 K.

68
Yield Strength, ay (MPa)

Figure 10. Inverse relationship between fracture toughness and yield

stress for austenitic stainless steels at 4 K.

69
Relative

Amount

of

BCC

Martensite

Figure 11. Relative amounts of a' (bcc) martensite formed during

tests of compact specimens at 4 K. (Arbitrary Units.)

70
STRESS,

cr

(GPa)

0 10 20 30 40 50

ENGINEERING STRAIN, e (%)

60

250

200

150

100

50

70

Figure 12. Engineering stress-strain curves for Fe-18Cr-10Ni steel

heat no. 1, representati ve of low interstitial C+N contents .

71

STRESS,

CT

(ksi)
STRESS,

er

(GPa)

Figure 13, Engineering stress-strai n curves for Fe- 1 3Cr- 1 ON i steel

heat no, 4, representati ve of medium interstitial C+N

contents

.

72

STRESS,

cr

(ksi)
STRESS,

cr

(GPa)

Figure 14. Engineering stress-strai n curves for Fe- 1 8Cr- 1 ON i steel

heat no. 8, representati ve of high interstitial C+N

contents

.

73

STRESS,

(T

(ksi)
TRUE

STRESS,

cr(GPa)

Figure 15. True stress-strai n curves for Fe-19Cr-10Ni steel heat

no. 1 (serrations at 4K are not shown).

74

TRUE

STRESS,

a(ksi)
TRUE

STRESS,

cr(GPa)

2.4

Fe-18Cr-10Ni

Heat No. 4

Figure 16. True stress strain curves for Fe-18Cr-10Ni steel heat

no. 4 (serrations at 4K are not shown).

75

TRUE

STRESS,

CT(ksi)
TRUE

STRESS,

a(GPa)

Figure 17, True stress-strain curves for Fe- 1 8Cr- 1 ON i steel heat

no. 8 (serrations at 4K are not shown).

75

TRUE

STRESS,

cr(ksi)
TENSILE AND FRACTURE PROPERTIES OF MANGANESE-MODIFIED

Fe-18Cr-3Mi TYPE AUSTENITIC STAINLESS STEELS

R. L. Tobler and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

A series of ten low-carbon AISI -304-type austenitic stainless

steels having 1 to 6% Mn (by weight) and 0.1 to 0.2% N were produced

and tested to determine the effect of these elements on properties at

4 K. Tensile tests (at 295, 76, and 4 K) and J-integral fracture

toughness tests (at 4 K) were conducted on developmental steels containing 18.25 to 19.50% Cr, 7.9 to 8.75% Ni, and 0.02 to 0.03% C. All

steels were hot-rolled at 1450 K (2150° F) from ingots to 25.4-mm (1-

inch) plates. The 4-K yield strengths ranged from 620 MPa (90 ksi) to

1068 MPa (155 ksi), increasing strongly with nitrogen content. Unacceptably low toughness was observed in the low manganese compositions,

but the fracture toughnesses of alloys containing 6% Mn were equivalent

to those of conventional AISI 304 stainless steels.

77

INTRODUCTION

The AI S I -300 series stainless steels are austenitic Fe-Cr-Ni alloys

offering relatively low strength but excellent cryogenic ductility and

toughness. Recently, nitrogen-strengthened grades, such as AISI 304 N or

AISI 304 LN, have attracted attention as possible substitutes for applications demanding higher strength. Nitrogen is a relatively inexpensive

strengthener, and it stabilizes the austenitic structure with respect to

martensitic phase transformation. A disadvantage of the nitrogenated

grades is that they are more difficult to produce and fabricate. To

overcome this disadvantage, manganese additions have been recommended.

The potential advantages of manganese-modified type AISI 304 LN steels

are (1) consistent fabricabil ity , (2) consistent weldability, and (3) reliable supply in all product forms [1]. It was the purpose of this

study to investigate the effects of manganese on the 4-K mechanical

properties of nitrogenated AISI 304 L type alloys. Therefore, a series

of steels containing 1 to 6% Mn and 0.1 to 0.2% N were prepared, tested,

and compared with previous results for Fe-18Cr-10Ni stainless steels [2,3].

MATERIALS

Ten stainless steel plates were obtained from a commercial steel

manufacturer . These steels had chemical compositions as listed in

table 1 and grain size and hardness values as listed in table 2.

All of these steels were produced from 225 N ingots (12.7 x 12.7 x

17.8 cm) which were hot-rolled at 1450 K (2150° F) to plates approximately 25-mm thick. The plates were annealed at 1340 K (1950° F) for

one-half hour, then water quenched and acid pickled. After trimming,

the average plate size was 25.4 x 140 x 400 mm.

79
Figure 1 shows one example of as-received plate. These steels were

free from corner checks (edge tears), which sometimes may occur in

nitrogen-alloyed austenitic stainless steels. Figure 2 shows an example

of corner checks as they appeared in an Fe-18Cr-10Ni-0.03C-0.024N-l.5Mn

stainless steel of a previous study [2], In the previous study a

series of nine Fe-18Cr-10Ni-l .5Mn stainless steels were hot-rolled at

1560 K (2350° F) from ingots to 25.4-mm-thick plates, and corner checks

were observed in steels having carbon plus nitrogen (C + N) contents of

0.270 wt.% or greater.

EXPERIMENTAL PROCEDURE

Tensi 1

e

Tensile tests were conducted using the cryostat described by

Reed [4]. Low temperatures were achieved by immersing the specimen and

cryostat in liquid nitrogen at 76 K and liquid helium at 4 K. The

specimen geometry is shown in fig. 3. The specimens were machined in

the transverse orientation. The load-vs-deflection curves were recorded

using the outputs from a commercial load cell and strain-gauge extensometer. The specimen yield stress at 0.2% plastic strain, ultimate tensile stress, elongation, and reduction of area were measured. Estimated

strain sensitivity is 50 ym at 4 K, and low-temperature accuracy is

estimated to be ± 500 ym at 4 K.

Fracture Toughness

J-integral tests were necessary because austenitic stainless steels

typically fail in a ductile (elastic-plastic) mode with gradual crack

extension at 4 K. Therefore, single-specimen J-integral tests were used

to measure fracture toughness. The technique incorporates a computer to

enable digital data acquisition and J-resistance curve plotting during

80
the tests, as described previously [2,3]. The fracture toughness

parameter measured in these tests is Jj

c

, the critical value of the Jintegral at the beginning of crack extension, which is a material

property enabling quantitative comparisons of ductile materials. Jj

also enables an estimate to be made of Kj , which is the plane-strain

fracture toughness as conventionally measured by the ASTM E 399 test

method.

The compact fracture toughness specimens shown in fig. 3 had a

24.6-mm thickness (B) and 50.8-mm width (W). The notch was machined in

the TL orientation, as defined in ASTM E-399-78a [5].

The fracture specimens were fatigue cracked at 76 K to a crack

length, a, between 27.25 and 31.44 mm (g - 0.535 to 0.620). The maximum

load used in fatigue precracking was 22 kN. The specimens were then

loaded in displacement control. As the load increased, partial unloadings

were performed periodically. The load reduction during unloading was

typically ten percent of the maximum load. Using the crack-length-vscompliance correlation, the crack length at each unloading was inferred

and used to obtain the crack extension increment (Aa) from the difference

of the initial precrack length. The J value at each unloading point was

calculated using the expression: J = -^ where A is the Merkl e-Corten

factor, A is the area under the load-vs-defection curve at the unloading

point, and b is the specimen ligament (b = W - a).

J-resistance curves consisting of many sets of J and Aa points were

generated. The Jj

c value for each steel was taken at the intersection

of the J-resistance curve with the blunting line. The blunting line is

given by J = 2a^A

a where is the flow stress, which is equal to the

81
average of the yield and ultimate tensile strengths. It was discovered

that the final crack extensions, Aa, measured on the fracture specimens,

were 20 to 200% greater than the values inferred by the computer. The

error between actual and inferred crack extensions increases as Aa

increases, and is lower at the Jj measurement point. Owing to this

bias, the Jj values reported here are estimated to be 15% higher than

would be expected from multiple specimen tests. Scatter in the Jresistance curves led to uncertainties of ± 15% in the values of J T . Ic

The validity of J, results was determined by the size criterion,

which requires that

B, b, a

Previous results have shown that valid results at 295 K are impossible

for annealed austenitic stainless steels of 25-mm thickness, so a complete temperature dependence of Jj was not attempted. Only 4-K values

are reported.

The 4-K Jj values were converted to Kj estimates, denoted Kj

c (J),

using the following equation:

K

Ic<

J >

1/2

where E is Young's modulus and v is Poisson's ratio. The values of E

and v at 4 K were estimated as 206.8 GPa and 0.30, respectively.

RESULTS AND DISCUSSION

Tensile Properties

The tensile properties of the steels tested in this study are

listed in table 3 and plotted in figs. 4 through 6.

As shown in fig. 4, the yield strength (o ) data for the present

Fe-18.25 to 1 9. 5Cr-7. 90 to 8 . 7 5 N i steels (open symbols) agree well with

previous data trends for Fe-18Cr-10Ni steels (solid lines). The yield

82
strength is highly sensitive to interstitial concentration, so differences in other metallurgical parameters (such as processing temperature,

grain size, and nickel content) can be neglected for this comparison.

The 295-K data confirm a previous conclusion: the yield strength is a

linear function of C+N content [2], The 4-K data also confirm the

existence of a kink in the a -versus-C+N plot at C+N contents between

y

0.127 and 0.157 wt.%. In the previous study, this kink was attributed

to the structural instability of the austenitic phase at these interstitial contents [2]. In the present study there is a clear trend of

increasing a values with higher Mn contents at C+N contents from 0.131

to 0.138 wt.%. This effect is probably due to the stabilizing effect of

Mn additions at this critical interstitial content.

The ultimate tensile strength results of the present study, shown

in fig. 5, are also in general agreement with expected trends. The data

at room temperature fall on the same trend line that was previously

reported for the Fe-18Cr-10Ni stainless steels. The 4-K data show a

minor variation owing to the austenite stabilizing effect of Mn modifications. As more manganese is added at a given interstitial concentration level, the austenite is progressively more stable during

tests at 4 K; as a result, less martensite is formed during testing and

ultimate tensile strength is reduced as Mn concentration is increased to

6%.

Figure 6 shows the spread of the elongation and reduction of area

measurements for the ten steels. Not all of the ductility values are

typical of commercial stainless steels. In particular, the reduction of

area values for the developmental steels containing 1 to 4% Mn are lower

compared with values previously reported for Fe-18Cr-10Ni steels [2] or

33
commercially available AISI 304 LN steels. Only the 6% Mn-modified

steels possess ductility values at the expected levels.

Fracture Toughness

The Jj results obtained for one specimen of each of the ten different steel compositions tested are reported in table 4, and the corresponding Kj (J) values are plotted in fig. 7.

As shown in fig. 7, the Kj (J) estimates follow regular trends

depending on Mn and N content. Solid lines have been drawn to connect

data representing constant Mn contents; for each Mn content the 0.1 N

composition is represented by the symbol at the lowest yield strength,

whereas the 0.2N composition is represented by the higher yield strength,

and 0.15N is intermediate. From fig. 7 it is clear that higher N at a

given Mn content lowers the Kj (J) value. It is also clear that Mn has

a favorable effect, raising both the strength and the toughness of these

developmental alloys.

The major observation from fig. 7 is that the 6%-Mn-modified steels

have acceptable fracture toughness, whereas the modifications containing

1 to 4% Mn display inferior properties. The 1 and 2 % Mn steels are

1/2 particularly poor, having estimated Kj

c values as low as 117 MPa*m .

Their load-deflection fracture toughness test records showed a nearly

linear-elastic behavior, with regions of fast crack advance (pop-ins).

This evaluation is based on comparison of the present results with the

previous results for a series of nine Fe- 1 8Cr- 1 ONi stainless steels containing variable C+N contents (for all of these steels the same Jj

c test

procedures were used). As shown in fig. 7, the 6% Mn-modified steels

exhibit an inverse K

y (J)/ct relationship at 4 K and, within the measurement error indicated by the error bars shown in the figure, the

34
inverse-linear trend falls along the lower bound of the scatter band of

data for the Fe-18Cr-10Ni steels.

The cause of the unexpected, low Kj (J) values of the 1 to 4% Mn

grades is not clear, but it may be related to chemical composition or

chemical segregation. The nominal compositional variations between the

1 to 2 i Mn steels of this study and the Fe-18Cr-10Ni-l .5Mn steels of the

previous study include a 1.25 to 2.5% difference in nickel content.

Metallurgical factors, such as interstitial segregation, may account for

the low fracture toughness of the 1 to 4 % Mn steels of the present

study. However, light microscopy revealed no evidence of sensitization

or other obvious microstructural differences.

Evidence for interstitial (nitrogen) segregation to grain boundaries

in the 1, 2, and 4% Mn steels was obtained by scanning electron microscopy performed on the fracture surfaces of the compact specimens subsequent to Jj testing at 4 K. Figure 8 presents typical results. The

specimens containing 4% Mn or less exhibited a significant portion of

brittle i ntercrystall i ne fracture, whereas the three heats containing

6 % Mn (heats 72, 76, 78) all showed a more completely dimpled fracture

mode. The lower manganese heats (69, 70, 73, and 74) also show more

intercrystalline fracture than was observed in the Fe- 1 8Cr- 1 ONi steels

of our previous study. The only apparent distinction between the

thermal -mechanical treatments of the two alloy series was the temperature of hot-working of the billet. The hot-rolling temperature of the

initial 304 LN alloy series was 1560 K (2350° F) and the rolling temperature of the 1-6 wt.% Mn alloy series was 1450 K (2150° F). Perhaps

this higher deformation temperature is required to preclude excessive

85
nitrogen segregation to grain boundaries. Further research is continuing to identify the distinctions between the two alloy series.

CONCLUSION

The data of this study show that 6 % Mn-modified AISI 304 LN stainless steel can be produced, having conventional tensile and fracture

toughness properties equivalent to standard AISI 304 LN stainless steel

when tested at 4 K. The Mn-modified stainless steel offers potential

processing and fabrication advantages when compared with the standard

grade, which contains 1 to 2% Mn.

ACKNOWLEDGMENTS

The stainless steel plates were produced by Armco, under the direction of R. H. Espy of Armco 1 s Research and Technology division. Dr.

E. L. Brown of NBS contributed valuable metallographic assistance and

technical advice, and D. Baxter of NBS conducted many of the tensile and

fracture toughness tests. This assistance is gratefully appreciated.

REFERENCES

[1] Espy, R. H., "Stainless steel alloy development," presented at

NBS/DoE Workshop on Materials at Low Temperatures , Vail, CO,

October 16-18, 1979.

[2] Tobler, R. L., and Reed, R. P., "Interstitial carbon and nitrogen

effects on the tensile and fracture parameters of AISI 304 stainless steels," in Material s Studies for Magnetic Fusion Energy

Applications at Low Temperatures-III . NBSIR 80-1627, R. P. Reed,

ed.. National Bureau of Standards, Boulder, CO, June 1980, pp. 17-48.

86
[3] Tobler, R. L., Read, D. T., and Reed, R. P., "Strength and toughness relationship for interstitial ly strengthened AISI 304 stainless steels at 4 K," in Materials Studies for Magneti c Fusion

Energy Applications at Low Temperatures-IV . NBSIR 81-1645, R. P. Reed

and N. J. Simon, eds.. National Bureau of Standards, Boulder, CO,

June 1 981 , pp. 37-76.

[4] R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range

300 to 4 K," in Advances in Cryogenic Engineering , Vol . 7, Plenum

Press, NY (1961), pp. 448-454.

[5] Standard Method of Test for Plane Strain Fracture Toughness of

Metallic Materials (Designation E399-74) , in 1 974 Annual Book of

ASTM Standards , Part 10, American Society for Testing and Materials, Philadelphia, PA (1974), pp. 432-451.

87
Table

1.

Compositions

of

modified

type

304

LN

stainless

steels.

O r— o r— o o 1o CD o CD o CD o r— LO LO LO LO | r— 1

11

r— p— 1CM 1CM r— <XJ (CXI CM 1

1(1

22 o o o o o o o o o o o o o o o o o o o o o o o 00 o cd o o o 00 o o o CD o l o o o o 3 CO co co CM co CM co co co CM CO CO co CM oo oo oo co CO CO co o o o o o o o o o o o o o o o o o o o o o CM o , o t o CO o CM o CM o l o ( o l o ( o CO co CO CO CO CO co co CO co CO co CO CO CO CO CO co CO CO o o o o o o o o o o o o o o o o o o o o

LD CO LO co LO o LO CM o o CD o o CO CM CD

•i— CO co CO CT> 00 CD 00 CD CO CD 00 oo CM co CM CO 00 00 00 CO CO 00 oo r^. r- i"--. CO 00 00 CO

LO CD LO o LO CM LO CD o co o LO o LO o 00 o o CM

CXI r— C\J f— CM CM CM o LO co LO CO LO CM LO ^3" CD CD

s- o CO 00 00 00 CO CO 00 00 cd CD CD CD CD CD CD CD CO CO 00 CO 1 “ - r—

r r— 1 1 1 1 1 1 o CO o co o CD o o o o o LO o o CM o CXI •1— LO LO LO CO LO LO LO co LO co LO co LO co LO co LO co LO co oo o o o o o o o o o o o o o o o o o o o o o CO o co o CO o ^3* o LO o CO o o cj- o o ^3" r— 1 r— r— 1 f—

r— r—-f— [ 1

f— r— r— r— 1 1— 1 f— p—

00 o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o C\J o - o CO o CO o CO o oo o CM o o CO o ^3"

C\J C\J C\J CM CM CM CM CM CM CM CM CM CM CM CM CXI CM CM CM CM

Q_ o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o

C\J o o o CD o o CD o r— o LO o CD o oo o o C o o o o o CO o CO o ,

o o o CO o o o o CO - - CM CM CO CO LO - r— CM CM ^3* CO co LO ^3" co LO o LO o 00 o LO o o LO o CM o o co o CO o r—

C\J CM CM CM C\J CM CVI CM CM CM CM CM CM CM CM oo CM CM CM CO o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o o

CL) -u • O o CD o f— CM CO ^3* LO CO 00

<_> CO

4-> CTi o CM CO ^3" LO CO oo

03 • co r^. O) O o o o o o o o o o o

zc 22 oo 00 CO 00 00 CO co 00 CO CO F

cu • -M O

1—1

21 1

CM co LO co 00 CD o

r—

03 03 03 03 03 03 03 03 03 03 3 3 3 3 3 3 3 23 D o e -t-> e 4-> E +-> E +J E 4-> E +J E +-> E +-> E 4-> E 4-> •r— o o <_> °r“ u •1— u •r— CJ •r— u •r- o •r (_> •r— o

03 03 03 03 03 03 03 ro 03 03 03 03 03 03 03 03 03 03 03 03

88
Table 2. Grain size and hardness (Rockwell B) of tested stainless steels 3

.

Code Number Carbon+Ni trogen

(wt. %)

Hardness

( r

b

)

Av. Grain Size

ASTM No. (mm)

69 0.135 82 5.4 (57)

70 0.138 88 5.5 (55)

71 0.135 84 5.8 (49)

72 0.131 84 5.4 (57)

73 0.215 88 5.5 (55)

74 0.212 94 5.5 (55)

75 0.214 90 5.6 (53)

76 0.243 90 5.6 (53)

77 0.176 89 5.6 (53)

78 0.181 90 5.6 (53)

a Grain size was measured on planes normal to the transverse

direction of the plate stock.

39
c

o

•f- rd 4-> <D ^-n LO CD ,

LO oo LD LD CD CNJ LO o ,

00 o cn CO , CO CNJ oo LO cn , cn cn CO o

o S- &S • Z3 <C - CNJ cn oo oo OvJ i

cn O 1 * CD CD co o CO co cn CNI CNI i

CNI CNJ o r— r— CNJ r-N T3 CD OO CVJ LD CNJ *3

* LO CO LO LO co LO co LO 1 LO CO CD LO CNI LO LO co LO LO CO

OJ 4-

CC O

co

+J

nga- (%) r- LO CO CD CNJ Cn CD CD LO CO co LO CNJ I-N LO 00 •— CD CD CO LO O CO CNJ O rCD , o LD 00 O f^N Cn CNJ IN. CO CO CO CNJ , , LO 00 CNJ CO CO CO CNJ 00 CO , o CD LO OO CD LO CNJ LO CO CD LO LO CO LO co CNJ LO CNJ co CD LO CO CD LD co LD LO *3

*

<D

S-d

4

-

> cd

5-

O)CLE

CL>

O)CL

CO

O)

+->

OJc

o

CO

OJ

+-> S-<uQ.o

S-o.

ai

CO

c:

CD

co

<D -O

ra

ai -c

-(-> -M

ra cn E c

•i- cl) +-> s_ r— -4-> => (D

+->CD

D fO

3 a.

CD

S-

+-> go

OJ

ra

>•, 0-

CL) Sra

s_

cdCLE

cd

CO

CD

4-> C

aj r—

+-> <o c c O -I—

<-j e

o

CO CO CD CD CO I - CNI r- C\J i

id in rv co ld cd

CD >3- .

00 O CO ID LO

CO CD CD

i— CNJ o CD CO ID

cn cni co

LO CO o CO CD CO

CNJ ON «^|- 03 ON CNJ CO LO LO

LO co

«d- i-n co

CD «^r CD

co cn «d- cn cnj oo CO -=3- LOCNJ LO CNJ CO

co oo lo in own

CO CO CO CO ID

CO LD CO 00 I

t"~ co .— cn i— cd

CNJ LO CO CNJ LO CD

CNJ CO CNJ

in. cni cni cnj lo i-n

LO O LO

LO 00 LO

cnj in in

N 'J- fO O O 00

co oo co

LO CD CO

cnj cn lo

co in. cn

cn i— co

CNJ 00 CO co r^. cn

co cn co l— LO

co oo o

CD N i— O i—

CO O LD CNJ I— O co i . oo co i-n cn

inco«jcn i

CNI

lO CD dcn i

CNI

in id ct

cn i-n- CNJ incodcn r''. CNI

in id cf

cn in.

CNJ

in co -d

cn in

CNJ

in co dcn r CNI iniodcn r-N

CNI

in co d in co

cn in. cn i

CNI CNI

o

o

o

o

o

o

o

o

oCNI o o o LO LO

CNJ CNJ CNJ i

r—

• • • • • o o o o o

+-> c

aj

+-> r- c 03 o sc O -r- CNJ^fCD,— CNJ «=J- CO CO

C

OJa

ajCL

oo

cn

CD O r—

i— IN.

CNI IN

COIN

•nT IN D CD IN. IN IN IN. 00

[N

90

*To

convert

from

MPa

to

ksi

,

divide

by

6.895.
Table

4.

Yield

strength,

flow

strength,

and

fracture

toughness

at4

K.

i

i— C\J cn cnj

«d- co

co LD o o

CM cm

-O

CM LD o o o o o o LT> o o

E • • o CO 01 PO o o 00 LO CO LO

1—

H

r~3 o CO o CO co co cn i

00 CO n JXL r— '— cm C\J «— r— '—

PO CO 00 1— PO 00 CM p-» r". CM PO LO i— r— r— C\J PO PO

fd

LO o 00 LO r-~ CM r— LO PO

PO CNJ PO PO CNJ CM

r— r— i

ii— i

CO r— PO LO PO CO CO PO r-^ i

f— CO CM LO 00 LO 00 LO LO o CO CO 1^ 00 cn cn O co cn

cr<yE

o

cuQ.

C/1

cn o i— c\j co r-* r^. p-. po LD CO

p-» r-^ r-. r-»

«d- co

p-^ co

r».

91

To

convert

from

MPa

to

ksi,

divide

by

6.895.

1

2

2

To

convert

from

kj/m

to

in-lb/in

,

multiply

by

5.709.

To

convert

from

MPa*/m

to

ksi*/i"n,

divide

by

1.099.
LIST OF FIGURES

Figure 1

.

Figure 2.

Figure 3.

Figure 4.

Figure 5.

Figure 6.

Figure 7.

Figure 8.

As-received 25 mm-thick stainless steel plate (heat 76,

Fe-1 9. 50-7 . 8 3 N

i - 5 .8 Mn), showing freedom from corner

checks

.

Corner checks as they appear in hot-rolled nitrogen-modified

Fe-1 8Cr- 1 ON i stainless steel of a previous study (see text).

Compact fracture and tensile specimens used in this study.

Effect of carbon plus nitrogen interstitials on yield strength

of stainless steels, with temperature as a parameter.

Effect of carbon plus nitrogen interstitials on ultimate strength

at two test temperatures.

Ductility of stainless steels.

Fracture toughness-versus-yield strength results for Mn-N

modified steels of the present study, as compared with previous

results for AISI 304 type stainless steels.

Scanning electron micrographs of 4 K fracture surfaces of low

toughness (heat 69, top) and high toughness (heat 72, bottom)

austenitic stainless steels.

92
LO

r"» S4-> CD

fC

C

CD 5- —

SZ

O 

o

CO E

CD o

+•> s_

CO 4- i

Q_ Eo i

-o

CD CD CD CD •M s_ CO 4—

CO cr> CO c.

CD •i— i

S C o

•i— sz

ro CO

+-> CO PI c u S •p— JC CO

+-> • 1 LO E 1 E •r— LO OO

CNJ CO •a

CD i > %-

•1

CO

CD LO • CJ • CO

CD cr> oc

i. i

u | i CD

co CD SI < Li. u . i

CD

S_3

CD • i

U_

93
94
COMPACT SPECIMEN

Thickness = 24.5 mm

Thickness = 2.54 mm

Figure 3. Compact fracture and tensile specimens used in this study.

95
YIELD

STRENGTH,

crv

(MPa)

Figure 4. Effect of carbon plus nitrogen interstitials on yield strength

of stainless steels, with temperature as a parameter.

95
ULTIMATE

TENSILE

STRENGTH,

cru

(

MPa)

CARBON + NITROGEN (wt.%)

Figure 5. Effect of carbon plus nitrogen interstitials on ultimate strength

at two test temperatures

.

97
ELONGATION

(%)

REDUCTION

OF

AREA

(%)

Figure 6. Ductility of stainless steels.

93
YIELD STRENGTH, cr

y

(MPa)

Figure 7. Fracture toughness-versus-yiel d strength results for Mn-N

modified steels of the present study, as compared with previous

results for AISI 304 type stainless steels.

99
(a \i 100 ym

(b) 1 00 ym

Figure 8. Scanning electron micrographs of 4 K fracture surfaces of low

toughness (heat 69, top) and high toughness (heat 72, bottom)

austenitic stainless steels.

1 00
INTERSTITIAL CARBON AND NITROGEN EFFECTS ON THE

CRYOGENIC FATIGUE CRACK GROWTH OF AISI 304 TYPE STAINLESS STEELS

R. L. Tobler and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Constant-load-amplitude fatigue crack propagation (FCP) rate measurements are reported for AISI 304 type stainless steels having variable

carbon-pl us-nitrogen (C+N) contents. Rates at stress-i ntensi ty factors

1/2 from 20 to 70 MPa*m were measured using 25-mm-thick compact specimens. The Fe- 1 8Cr- 1 ON i steels tested exhibited partial martensitic

phase transformations during tests at 76 and 4 K, but not during tests

at 295 K. At high C+N contents (0.187 wt.% or greater), the FCP resistance was lower at 4 K than at 295 K. At low C+N contents (0.067 wt.%),

the FCP resistance at cryogenic temperature was significantly higher

than at room temperature. The improvement in steels of low C+N contents

was associated with a transition from partially intergranular failure at

295 K to completely transgranular failure at 76 and 4 K.

101

INTRODUCTION

When James [1] reviewed the subject of fatigue crack propagation in

austenitic stainless steels in 1976, few cryogenic data were available.

Since then, fatigue crack propagation rate measurements (da/dN) at

intermediate stress intensity factor ranges (aK) have shown that stable

( nontransforming ) austenitic stainless steels display improved resistance to fatigue cracking at extreme cryogenic temperatures, whereas

metastable grades are less predictable [2-5]. A metastable stainless

steel is one in which the austenitic phase in response to stress or

plastic deformation at low temperatures transforms to some degree to

body centered cubic (a

1 ) or hexagonal (e) martensitic phases.

Metastable AISI 304 stainless steel (Fe-18 to 20Cr-8 to 12Ni) is a

common cryogenic structural alloy, whose mechanical properties are

highly sensitive to composition; a broad range of mechanical property

combinations can be achieved by varying the C+N contents, because these

elements significantly stabilize and strengthen the austenitic phase. A

previous study described the effects of C+N contents on the 4-K fracture

toughness properties of a series of Fe-18Cr-10Ni developmental stainless

steels [6]. In the present study, the same series of steels was tested

to evaluate fatigue crack propagation behavior.

MATERIALS

Nine stainless steel alloy plates with an Fe- 1 8Cr- 1 ON i base were

tested. These plates had a nominal base composition falling within the

limits set by ASTM specification A 240 for AISI 304 stainless steel.

The C+N levels varied, however, with carbon at 0.03, 0.06, or 0.09 wt.%,

and nitrogen at 0.04, 0.12, or 0.24 wt.%. The nine plates were produced

103
from three 1 36- kg (300-lb) vacuum-i nduction-mel ted heats, split with

respect to carbon level, and teemed into 76 x 200 x 360 cm hot-topped

cast-iron ingot molds. The ingots were then reheated and soaked at

1561 K (2350° F), hot-rolled to 25 .4-mm-thick plates, and air cooled.

The plates were finally annealed at 1332 ± 7 K (1937° ± 13° F) for 1 h

and water quenched. The mill chemical analyses, hardnesses, and grain

size measurements are listed in Table I. Tensile properties are listed

in Tabl e II.

SPECIMEN

The modified compact specimen geometry of Figure 1 was used (the

basic geometry is described in ASTM Method E-399-74 [7]). The modification consisted of mounting knife edges for clip-gage deflection (6)

measurements at the load line. The specimen thickness (B) was 25 mm,

and the width-to-thickness ratio (W/B) was 2. The notch was machined in

the TL orientation [7], and the crack propagated in a direction parallel

to the plate rolling direction.

PROCEDURES

The fatigue tests were performed using a 1 00- kN servo-hydraul ic

testing machine and cryostat [8]. Room-temperature tests were conducted

in ambient air (295 K) at a relative humidity of 10 to 20%. Cryogenic

tests were performed with the specimens submerged in a stainless steel

Dewar containing liquid nitrogen (76 K) or liquid helium (4 K)

.

Specimens tested at 295 and 76 K were precracked at the same temperatures, whereas specimens tested at 4 K were precracked at 76 K.

During precracking at 76 K, the number of cycles to initiate a 1.9-mm

crack was recorded. The machine was operated in load control, using a
sinusoidal load cycle at frequencies between 20 and 25 Hz. The load, P,

was accurate to ± 3%. The load ratio, R (Pm

,-„/Pm3W ) > was constant at

nil rl maX

0.1. The maximum load for a given test was held constant at a value

between 22 and 24 kN.

Crack growth-to-a/W ratios of about 0.65 were monitored by specimen

compliance (6/P) measurements. The compliance measurements were correlated

with the average crack lengths based on measurements at the specimen

edges, quarter points, and center of thickness [5,8], Crack curvature

was such that the average crack length was 93% of the center crack

length for a/W ^ 0.6. Robert's crack length-compliance solution was

corrected (shifted a few percent to agree with the experimental measurements

for the final crack length for each test specimen) and then used to

infer all other crack lengths from compliance readings taken during

fatigue tests [5,8]. A computer was used to fit the a-versus-N curves

with third order polynomial expressions, to differentiate the curves to

obtain da/dN, and to calculate the correspond!' ng stress intensity factor

ranges, aK:

AK = K -K . = (P -P . )B~V1 / 2 [f(a/W)] (1) max min max min L w /J v '

where f ( a/W ) is given by Roberts [9]. Data were obtained from one

specimen per material per temperature.

Magnetic measurements verified that cyclic loading of the Fe-18Cr1 ON i steels during crack growth experiments at 76 K and 4 K caused

formation of a

1 martensite. The a' martensite is ferromagnetic, whereas

the original austenite and e martensite are not. Therefore, by using a

torsion bar magnetometer [6] on the compact specimen fatigue surfaces

105
subsequent to testing, the amount of a' martensitic transformation was

determi ned.

RESULTS

Crack initiation at 76 K

The number of load cycles required to produce a 1.9-mm crack at

76 K for the particular specimen geometry and loading conditions used in

the present tests are plotted in Figure 2. As C+N contents are increased, the resistance to crack initiation decreases rapidly at first,

but reaches a plateau at C+N = 0.157 or greater. The effect is large,

amounting to a five-fold variation of initiation time.

Crack propagation

For the stress intensity factor ranges studied, the FCP rates for

Fe- 1 8Cr-l ON i stainless steels conformed to power law equations of the

form

:

da/dN = C( AK)

n (2)

where the parameters C and n are material and temperature dependent.

Evidence justifying the applicability of this equation to austenitic

stainless steels, and to AISI 304 in particular, is presented by James [1]

and Brose and Dowling [10]. In the present study, log-log plots of

da/dN versus AK were constructed to obtain the C and n values that

correspond to the ordinate intercepts at AK = 1 and the slopes, respectively

.

Table III lists the C and n values for all material and temperature

combi nations

.

A general presentation of results is shown in Figure 3. The roomtemperature data (Fig. 3a) for the present steels form a scatter band

106
that fits within the larger scatter band drawn by James [1] for many

commercial AISI 304 steels tested at different laboratories using ten

different specimen geometries and variations of test procedure. The

data for the Fe-18Cr-10Ni steel of our study having 0.067 wt.% C+N was

also found to be in close agreement with previous measurements for a

commercial AISI 304L steel of similar composition [5]. Therefore, our

results for laboratory heats are representati ve of commercial steels.

Figures 3b and 3c show the effect of temperature reductions on the

spread of present results. Reducing the temperature to 76 K improves

the FCP resistance of many of these steels. At 4 K, however, the high

C+N compositions are adversely affected. As indicated in Figure 3c, the

4-K scatter band encompasses the 295- K scatter band for the same steels.

The greater variation of fatigue crack growth rates at 4 K is attended

by significant differences in austenitic stability and tensile properties (Table II) for these alloys, owing to the effects of C+N at cryogenic temperatures.

Interstitial effects

Figure 4 presents the results for each heat at 295, 76, or 4 K.

Four heats were tested at 295 K. As shown in Figure 4a, the rates for

heats 1 and 3 are equivalent. The rates for heats 5 and 9 are also

equivalent, but lower compared with those of heats 1 and 3. Therefore,

higher C+N contents tend to improve fatigue propagation resistance at

295 K. Individual datum points are shown only for heats 1 and 9. The

specified C+N contents for these two steels (0.067 and 0.325 wt.%) represent the limits studied, and the indicated data spreads are typical.

107
Also shown in Figure 4a are results for an AISI 310 stainless steel

( Fe-25Cr-20Ni ) [5], which fall among the Fe-18Cr-10Ni results. At

295 K, all of these steels are stable with respect to the austeni te-tomartensite phase transformation, and all exhibit similar microfailure

modes. As shown in Figure 5, the failure modes at 295 K consist of

significant portions of intergranular failure. Scanning electron micrographs were taken from the fatigue surfaces where the aK value was

1/2 approximately 40 MPa*m . In each case, the failures show a number of

features, including secondary cracks, coarse transgranular failure

zones, and smooth, highly reflective grain-boundary facets. Occasionally, a sample would show delaminations. The exposed grain facets

frequently displayed slip-line markings, sometimes twins.

At 76 K (Figure 4b), the effect of C+N content on fatigue crack

growth rates is reversed from that at 295 K. The rates for heat 1 are

considerably lower than for heats 3, 6, 7, and 9, which are nearly

equivalent. Thus, the steel with the lowest C+N level has the lowest

rates at 76 K. At 76 K, partial martensitic transformations take place,

and transformation is most extensive for heat 1. Significantly, there

is also an obvious difference in microfracture mode for heat 1. As

shown in Figure 6, the 76-K fatigue surface reveals a completely transgranular microfailure, quite distinct from the partially intergranular

mode exhibited by this steel at 295 K. As a result of this transition

to a transgranul ar failure, the 76 K fatigue surfaces for heat 1 are

noticeably smoother than those at 295 K. In Figure 6, the tendency to

transgranular failure is also somewhat in evidence for heat 3. In fact,

as the C+N contents are increased, grain facets gradually become more

103
prominent on the 76-K fatigue surfaces. Heats 6, 7, and 9 at 76 K

exhibit the same degree of intergranular cracking observed in tests at

295 K.

At 4 K (Figure 4c), where all heats were tested, the results were

similar to those at 76 K. Again, heat 1 exhibits the lowest fatigue

crack growth rates. As C+N contents are increased, the 4-K rates increase rapidly at first, but the effect saturates at C+N ^ 0.128 or

more, and heats 4 through 9 have equivalent fatigue crack growth resistances at this temperature.

The relative amount of a' (bcc) martensite formed during fatigue

testing at 4 K is plotted as a function of interstitial content in

Figure 7. These data provide only an approximate indication of the

extent of phase transformation , because the magnetometer readings were

taken directly from the compact specimen fatigue surfaces, which are

unmachined. From Figure 7, it is clear that some martensite forms in

all of the alloys fatigued at 4 K. But more martensite forms in the

low-strength, low-C+N grades, because these grades are the least stable

and also have the largest crack-tip plastic-zone sizes.

Scanning electron microscopy results for specimens fatigued at 4 K

are shown in Figure 8. The findings closely parallel those at 76 K: at

low C+N contents, the failure mode is completely transgranul ar , but at

high C+N contents there is a partially i ntergranul ar failure mode.

Temperature effects

The results for three heats representing low, intermediate, and

high C+N contents are replotted in Figure 9, illustrating the temperature dependence of fatigue crack growth resistance.

109
Heat 1 (C+N = 0.067 wt.%) represents the lowest interstitial

content, lowest yield strength, and lowest austenite stability. For

this steel, temperature reductions to 76 or 4 K greatly improve the

da/dN values at a given AK value. The fatigue crack growth rates at 76

and 4 K are nearly equivalent and much lower than those at 295 K. The

rates for heat 1 are in excellent agreement with rates previously reported for a commercial AISI 304L and 310 stainless steels [5] at 76 and

4 K. As indicated previously in Figure 5, the improved fatigue crack

growth resistance at cryogenic temperatures is associated with a transition from partially intergranular failure (295 K) to transgranular

failure (76 and 4 K)

.

Heat 3 (C+N = 0.128 wt.%) represents intermediate interstitial

content, strength, and austenite stability. As shown in Figure 9b, the

rates at a given aK value decrease slightly between 295 and 76 K and

then increase between 76 and 4 K. Thus, temperature effects are minimized at intermediate interstitial content. This concurs with previous

reports of temperature-insensitive fatigue crack growth behavio in

commercial AISI 304 grade steel [5]. The scanning electron micros:: :y

results for heat 3 show a more well-defined intergranular failure mode

at 295 and 4 K, compared with that at 76 K.

Heat 9 (C+N = 0.325 wt.%) represents the highest interstitial

content, highest strength, and highest austenite stability. The effect

of a temperature reduction to 4 K on this heat is completely opposite to

the trend for heat 1, shown in Figure 9c. The 7 6- K rates are slightly

higher or nearly equivalent to the rates at 295 K, but the 4-K rates are

two or three times higher. The failure mode for heat 9 is partially intergranular at all test temperatures.

no
DISCUSSION

It was concluded from previous work that desirable fatigue crack

growth resistance at cryogenic temperatures may be achieved with alloys

having the austenitic structure and high elastic modulus [5]. Nevertheless, as shown in this paper, a broad range of properties may be displayed within a given austenitic steel family at 4 K, owing to compositional

variations

.

At room temperature, where the da/dN data represent crack propagation through s able austenite, increasing the C+N level tends to increase fatigue crack growth resistance, in accordance with previous

observations [2], although heat-to-heat variations are not very great.

The microfailure mode at room temperature always consists of a rough

transgranular surface plus a significant portion of intergranul ar

failure. The grain facets corresponding to the i ntergranular failure

are undoubtedly the reflective "cleavage-like" facets observed under low

magnification in previous studies of austenitic steels [5].

At cryogenic temperatures, where the strengths of some of the

steels were greatly increased and where martensitic transformations

occurred in varying degrees, the variation of fatigue-crack-growth

resistances was especially great. During tests of these steels at 76

and at 4 K, it can be assumed that the martensitic transformation occurred locally in the plastic zone ahead of the crack, so that transformation preceded material separation. The parameter measured in these

tests was the rate of fatigue crack propagation through partially martensitic material

.

At low C+N levels, the effects of martensitic transformation are

complex: strength is low, strain-hardening capability is high, and a

relatively large portion of martensite is formed. Since the martensite

111
transformation is an energy-absorbi ng process, a decrease of crack

propagation rates may result, as compared to rates under conditions

where the austenite is more stable.

At higher C+N contents, the strength level is high and strain

hardening capability is low. The amount of phase transformation is

reduced, and its effects may be masked by other factors, such as a

growing tendency to i ntergranul ar fracture. Intergranular fracture may

be encouraged at higher C+N contents if the grain boundary energy is

reduced by absorbed interstitials or if the grains are hardened relative

to the grain boundaries.

These results must be interpreted cautiously, because it is believed

that the cyclic stress-strain behavior, not the static tensile properties,

govern fatigue behavior [2]. In our work, the effects of C+N contents

and martensitic transformation on cyclic stress-strain properties were

not studied. The effect of the volumetric expansion accompanying martensitic

transformation is also unknown. The volumetric expansion in austenitic

stainless steels is about 3 %, and this should introduce residual stresses.

Separation of the temperature, transformation , strain hardening, and

residual stress effects is not easy and is beyond the scope of this

study.

A primary purpose of the present work is to compare the fatigue

crack growth rates of steels that may be selected for large cryogenic

machinery. Nine metastable Fe-18Cr-10Ni steels and an AISI 310 steel

are compared. The two alloys having superior fatigue crack growth

resistances are the fully stable (AISI 310) and the least stable

(AISI 304, heat 1). The 76- and 4-K rates of these two (Fig. 4) are

112
equivalent and less than the rates at 295 K. The equivalent performances at extreme cryogenic temperatures are surprising since the improvement of the metastable steel is associated with a failure mode

transition, whereas the improvement for AISI 310 must be attributed

solely to inherent temperature effects on material properties (i.e.,

there is no failure mode transition for AISI 310).

Fatigue crack propagation at 4 K may be an important design parameter for superconducting magnet structures in some magnetic fusion

energy devices. The rates reported here were determined primarily for

material characterization, but the results may also be useful for life

prediction, with precautions. Owing to the complex nature of the steels

tested, the data trends of Table III should not be extrapolated beyond

the range of actual data accumulation. The rates in the present study

were measured at R = 0.1, at constant load amplitude, and at intermediate aK. These conditions are not representative of many applications.

Before the fatigue life of a component can be accurately predicted, the

effects of stress ratio, variable amplitude loading, cycle frequency,

and deviations from the power law equation at threshold or very high AK

levels need to be evaluated. These effects have never been studied at

4 K.

SUMMARY AND CONCLUSION

A study was conducted to determine the effects of interstitial

C+N on fatigue crack propagation resistance of AISI 304 type stainless

steels at room and cryogenic temperatures. Rate measurements were

reported for conditions of constant load amplitude, at a stress ratio of

0.1, and at intermediate aK values, using compact specimens of the TL

113
orientation, taken from 25-mm-thick plate. The C+N content determines

the temperature dependence of fatigue crack propagation rates in these

al 1 oys

:

(1) At low C+N contents, the rates decrease at cryogenic temperatures, relative to those at 295 K,

(2) At intermediate C+N contents, the rates are relatively insensitive to temperature, and

(3) At high C+N contents, the rates increase at cryogenic temperatures, relative to those at 295 K.

The improvement for low C+N contents is associated with a failure mode

transition from partially intergranular failure at 295 K to transgranular failure at 76 and 4 K.

114
REFERENCES

1. James, L. A., "Fatigue crack propagation in austenitic stainless

steels," At. Energy Rev. 1_4 , 1 (1976), 37-85.

2. Pineau, A. G., and Pelloux, R. M., "Influence of strain-induced

martensite transformations on fatigue crack growth rates in stainless steels," Metal 1. Trans., 5^ (1974) 11 03-1 11 2.

3. Logsdon, W. A., Wells, J. M., and Kossowsky, R., "Fracture mechanics properties of austenitic stainless steels for advanced

cryogenic applications," in: Proceedings of the Second Inter -

national Conference on Mechanical Behavior of Materials , American

Society for Metals, Metals Park, Ohio (1976), 1283-1289.

4. Tobler, R. L. and Reed, R. P., "Fatigue crack growth rates of

structural alloys at 4 K," in Advances in Cryogenic Engineering ,

Vol . 22, K. D. Timmerhaus, R. P. Reed, and A. F. Clark, eds.,

Plenum, New York (1977), 35-46.

5. Tobler, R. L. and Reed, R. P., "Fatigue crack growth resistance of

structural alloys at cryogenic temperatures," in: Advances in

Cryogenic Engineering , Vol. 24, K. D. Timmerhaus, R. P. Reed, and

A. F. Clark, eds., Plenum, New York (1978), 82-90.

6. Tobler, R. L. and Reed, R. P., "Interstitial carbon and nitrogen

effects on the tensile and fracture parameters of AISI 304 stainless steels," in: Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures-III , NBSIR 80-1 627, R. P. Reed,

ed.. Fracture and Deformation Division, National Bureau of Standards, Boulder, Colorado (June 1980), 17-48.

115
7. Standard test method for plane-strain fracture toughness of metallic materials (designation ANSI/ASTM E399-78a)," in 1 979 An -

nual Book of ASTM Standards , Part 10, American Society for Testing

and Materials, Philadelphia (1979), 540-561.

8. Fowlkes, C. W. and Tobler, R. L., "Fracture testing and results for

a Ti-6A1-4V alloy at liquid helium temperature," Eng. Fract. Mech.

8 (1976), 487-500.

9. Roberts, E. J., Jr., "Elastic crack-edge displacements for the

compact tension specimen," Mater. Res. Stands., 9_, No. 2 (1969),

27.

10. Brose, W. R. and Dowling, N. E., "Size effects on the fatigue crack

growth rate of type 304 stainless steel," in: Elastic-Plastic

Fracture , ASTM STP 668, J. D. Landes, J. A. Begley, and G. A.

Clarke, eds., American Society for Testing and Materials, Philadelphia (1979), 720-735.

116
Table

I.

Stainless

steel

compositions

(wt.%),

hardnesses,

and

grain

sizes.

cdN

CD -r- CD <J~) fo E LO CO O LO LO 00 00 CM o

C D- co CD O 00 CD co CD

CD r— > 03 <C S- o

CQ

00 i—

OO 1— CO o co t— LO i—

CD <D c 5 CO co co CD CM CM CD CD r—

T3 r-^ r-- CO CO co CD

S~ O

ro O

3: a:

r^. r- co r*^ r^. <3- o r— LO

z: oo CD CM LO 00 i— r-^ CD CM + O o i— i— 1— CM CNJ CXI CO o o o o o o O O o O

CD o o

CO CM O = - '

= = CM = - o o O

00 co CD r^. O r— LO

C\J LO CO co LO CD CO LO co o o o O o o O O o o

o o O o o O o o o

o o 1

CD o

13 CM z z i

z z CM z z

C_) o o O

LO LO o o CD o CNJ CM ,—

21 o o o

o LO

on LO LO o o O

CD LO

CO o - = o - = o = = o o o

, LO

CNJ CM CM

Q_ O - = o = = O = = o o O

o CM e LO LO

21 • — — • — — • — —

r~* '

•— CD CD o CD CD

1

CO O CO O co CO o

CM

CD

Li— 03 03 03 03 03 03 03 03 03 _Q JO JO JO JO JO JO JO JO

O

r— CM CO LO LO C". CO CD

+-> fO

a> m

117
Table

II.

Tensile

properties

(one

test

per

temperature).

fO

cd

s_ c

“O

CD Cd

,— ,— LO LO CO LO

.— cm co (\j to to

r>. lo lo lo

co cm o ro co UTlTUD CM *3" LO LO *3" to LO

CT* LO 00 i— CM

cm lo *3

- lo co

LO LO LO O r— O r- ^ N*. LO LO CO

LO <3" LO LO

.— CT* CM LO LO

n* cr* r— co cr> ,

to lo <3- lo ro co

co

03 5^ Co

.— co 00 On lo

*— cr. cr> r— ro cm n ro cm 1— «=r co

cn n ^ lo «^r

cr* >^3- r\ cr* LO ^3- ro LO «=}-

LO CM co cr* LO CT* <3- cr* ro LO O- co LO O ^ O CM N

O LO CD LO r—

LO CO LO LO CO

CO O LO NON

CM CO CM N O CM

LO LO CO *d" CO CO

LU

LO (T> *T CO

cr* 1— lo o*> ct- LO -^3- LO «=j- LO ^ LO COCM 0> CM

LO ^3- LO

CM O O .— in n CO <3- LOCM ^3- O u~> cr* o LO ^3" N^

00 CO

LO CO

LO LO LO

CO LO LO

cr* co co

LO LO o co lo cr* o lo

<3- LO CO CM CO O

NIT) N N ID N

TD

CD

CM CT* 1— LO LO

CM "d- CM *3" 1

CM CM CO CM CO ^

CO CO O LO CTi CO

CM CO LO

cr* lo lo

N- 1— *3" CM LO N

CM 00 LO r- CO N ro co co ro n COcr* cr* lo

lo co co CO CO r—

co cm co lo cr* lo LD LD N N» .— CO ro cr* 1— co cr* cm

s-

<dCL^ LO LO «=f

cr* ncm

LO LO ^3-

CT> NCM

LO LO sj

cr* nCM

LOo

o

fN.

cr* o

o

o Ncr* CM CM o o

LOCM

COO

O

>>O 4->

<— 03 r— CD c m

CM CO <3- LO LO N^ CO cr*

118

*To

convert

from

MPa

to

ksi,

divide

by

6.895.
Power-law

fatigue

equation

[Eq.

2]

parameters

for

austenitic

stainless

steels

(R

=

0.1),

1/2

where

da/dN

is

in

mm/cycle

and

aK

is

in

MPa*m

'

.

c (XI o

r~ n—

cr, E

CU • CC a3 Q_

<1 '—

r-^ LO o o o *d- o O CO o o O CO o LO o O o LD LO LO •tf- LO LO co i-n ID ID LO LO LO LO LO CO LO CO CO O LO LO o

1 LO LO

1 1 O ID

1 1 LD LD LO LO

1 1 cj- «=a8 LD1 LO1 LO

CM CM (XI (XI CM C\J <XJ (XI CO CO C\J CM CM CM CM CM CM CM CM CO CM CM

o O i— r— o o o o ,

o o r— O i— O O r— r— 0 1— r— r— i

cn i— cn >— cn i— cn cn i— f— cn r— CT> ,

| | O 1 1 B 8 | 0 8 1 8 0 1 8 8 8 8 8 8 8 0 o o 1 o o o o o o o o o o o o o o o o o o o r— r— 1— r— i— r— f— r— t— r— r— f— f— r— f— p— f— r— f— p— r— o X

X X X XXX X X X X X X X X X X X X X X X

LO i— r-'. CM r-» 1— LD r— cn ,

1

Cj" i <— cn ,ID N ,— 1

CO CO CO CO M- •- CO 1— CO I— CO 1— r— CO CO r— CO ^3" 1— 1—

LD ' LO

cr> o cn cn r- co LO CM LO LO pn lo LO CM IN LD LtONS c

co «=f CO co CO CO CO CO ^i- CO CO CO co co CO CO CO D- CO CO CO

a)

S-D

+->

<T3 * LD LO *3" LO LO ^3* LO LO «=f lo ID LO Lt ID ID LD

S- CTi ct> r--. ct> pn CT> p-. cti cn r^. aj V CM CM CM CM CM CM

CLF

CL) 1—

a> -Q

-M

03

aj

c\j co LO lo co CT1 O

CO

GO

i—

<c

119

*To

convert

from

MPavfiPto

ksi/in,

divide

by

1.099.
LIST OF FIGURES

Figure 1. Compact specimen used for fatigue tests.

Figure 2. Crack initiation results for specific conditions at 76 K.

Figure 3. General comparison of da/dN vs. aK results.

Figure 4. Fatigue crack propagation behavior as a function of interstitial C+N contents with temperature as a parameter.

Figure 5. SEM views of the partially intergranular failure modes at

295 K.

Figure 6. SEM views of the fatigue failure modes at 76 K, showing the

transgranul ar (heat 1) to partially intergranular transition

( heat 9)

.

Figure 7. Relative amounts of a

1 (bcc) martensite formed during fatigue

tests at 4 K. (Arbitrary units)

Figure 8. SEM views of the fatigue failure modes at 4 K, showing the

transgranul ar to partially intergranular transition.

Figure 9. Fatigue crack propagation behavior as a function of temperature, with interstitial C+N contents as a parameter.

120
COMPACT SPECIMEN

Thickness = 24.5 mm

Figure 1. Compact specimen used for fatigue tests.

121
500

^ODJQ UllU0,

| D 94 DI 4 IUI 0| S3|0^00|{^|

122

0.10

0.20

0.30

Carbon

+

Nitrogen

(weight

percent)
FATIGUE

CRACK

GROWTH

RATE,

da/dN,

mm/cycle

Fe - l8Cr - lONi STAINLESS STEELS

Figure 3. General comparison of da/dN vs. AK results.

123
FATIGUE

CRACK

GROWTH

RATE,

da/dN,

mm

/cycle

Fe - 1 8 C r - 1 0 N i STAINLESS STEELS

Figure 4. Fatigue crack propagation behavior as a function of interstitial C+N contents with

temperature as a parameter.

124
Figure 5. SEN views of the partially i ntergranul ar failure nodes at 295 K.

125

heat

1

AISI

310
Mgure 6. SEM views of the fatigue failure modes at

76 K, showing the transgranul ar (heat 1) to

partially intergranul ar transition (heat 9).

126

40HM
Figure 7. Relative amounts of a' (bcc) martensite formed

during fatigue tests at 4 K. (Arbitrary units.)

127
LO

-P

rO

<v

jzr

Figure C. SEM views of the fatigue failure modes at 4 K,

showing the transgranul ar to partially inter-’

granular transition.

128

40HM
FATIGUE

CRACK

GROWTH

RATE,

da/dN,

mm/cycle

Fe -!8Cr — 1 0 N i STAINLESS STEELS

I

1 1—1—

1—1 III ^ i i

i

i i i i ^ TTTT

Heat 1 Heat 3 Heat 9

C + N= 0.067 wt.% C + N = 0.128 wt. % C + N =0.325 wt.%

: / /

' /:

E /

' /- // -

: / / / //

"

/ /

295 /

/

/ /

295 / /

and / / 4 / /

/ /

4

11 //

- // //

/ Z/

295

I/ 77

4

//

//

77/

i / i i 1 i i i i i i i 1 i i i i > i i 1 i i i i 20 40 60 80 20 40 60 80 20 40 60 80

STRESS INTENSITY FACTOR RANGE, AK, MPa • m

l/2

Figure 9. Fatigue crack propagation behavior as a function

of temperature, with interstitial C+N contents as a parameter.

129

ANOMALOUS YIELDING, TRANSFORMATION, AND RECOVERY EFFECTS IN

AISI 304L STAINLESS STEELS

R. L. Tobler, D. H. Beekman, and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Experiments describing the temperature dependence of yielding and

anomalous stress-strai n behaviors of metastable AISI 304L stainless

steels were performed. In these steels transformation to a

1 bcc martensite can be induced by thermal cycling or creep at stresses significantly lower than those required to form a' by simple monotonic

tensile loading. A recovery phenomenon is observed for tensile specimens unloaded and warmed to 295 K after plastic strain at 76 K.

1 31

INTRODUCTION

Austenitic stainless steels have a metastable face-centered cubic

structure (y). But at room temperature and below both body-centered

cubic (a 1 ) and hexagonal close-packed (e) structures form during cooling,

under applied elastic stress, or during plastic deformation. In the

commercial grades of Fe-18Cr-8Ni the martensitic phase transformations

may occur during service, depending on alloy content, service temperature, stress history, and metallurgical condition.

Various physical and mechanical property changes occur in conjunction with martensitic transformation in austenitic stainless steels.

The most significant consequences of transformation are volumetric

expansion (about 2%), magnetic transition from paramagnetic (y) to

ferromagnetic (a 1 ), and mechanical property changes.

In the Fe-Cr-Ni austenitic steels, the martensitic transformation

is believed to be responsible for anomalous temperature dependence of

the flow strength. Suzuki et al . [1] have suggested that the transition

to e is responsible for flow strength decreases at lower temperatures,

whereas Olsen and Azrin [2] have produced data that suggest a strong

correlation of the decline of flow strength at low temperatures with a

1

martensite. There are other ramifications of the role of martensitic

transformation in plastic deformation of austenite. This paper suggests

that thermal cycling and creep behavior are significantly influenced.

In particular, preliminary data are presented that suggest that austenitic alloys do creep at low temperatures and that thermal cycling (76-

295 K) serves to alter the creep rate.

133
EXPERIMENTAL PROCEDURE

The materials studied are AISI 304L stainless steels. One of these

steels had been previously tested [3], The second steel has the following composition in weight percent: Fe-18.3Cr-10.lNi-l.57Mn-0.021P0.019S-0.50Si-0.20Mo-0.21Cu-0.028C-0.039N. This steel was obtained as a

25-mm-thick plate, having an average grain diameter of 85 ym and a

Rockwell hardness of Rg74.

All tests were performed on dog-boned specimens 2.54 x 4.77 mm in

cross section, 50.8 mm in gage length, and transversely oriented with

respect to the plate rolling direction. Conventional tensile tests used

a universal test machine at a crosshead rate of 8.47 ym/s. Creep and

thermal cycling tests were conducted using a servo-hydraul ic test

machine in load control.

Clip gage extensometers and inked scribe lines were used to measure

strain. Test temperatures were achieved using liquid helium (4 K)

,

liquid nitrogen (76 K) , or a vapor cooling system [3].

In thermal cycling and stress tests, the specimens were loaded to a

specified stress at 295 K, cooled to 76 K in a Dewar of nitrogen, and

then warmed to 295 K in a jet of air after removing the Dewar. This

procedure was repeated 10 times in about 0.5 h.

The amount of body-centered cubic (a

1 ) martensite formed after

various tests was measured using a spring-loaded magnetometer. Five

magnetometer readings spaced across the gage length were taken on opposite sides of the specimen, and averaged. The accuracy of the a

1

martensite determination is estimated at ± 0.02%. The magnetic readings

had been previously associated with volume percent of the a

1 , a ferromagnetic phase [2].

134
RESULTS AND DISCUSSION

Tensile yield strength

Yield strength data for AISI 304L stainless steels are plotted as a

function of temperature in Fig. 1. A regular trend is expected for

stable austenitic stainless steels, whereas the data for AISI 304L

demonstrate an irregular temperature dependence: the yield strength

increases between 295 and 200 K, then decreases between 200 and 100 K,

and finally increases again between 100 and 4 K. The present results

follow the trend of Read, Reed, and Schramm [3], who tested a slightly

stronger heat (higher carbon content).

Owing to the anomalous behavior demonstrated in Fig. 1, the 76-K

yield strength of AISI 304L is only 40 MPa higher than at 295 K. For

the thermal cycling and creep experiments described below, this means

that for a given applied stress, a

Q

, there is not a great difference

between the applied stress to yield strength ratio (a /a ) at 295 and

a y

76 K.

Behavior during the initial stage of plastic deformation is illustrated by the stress-strain curves shown in Fig. 2 for the steel of

reference 3. During the first few percent of plastic strain, the

AISI 304L curves at different test temperatures cross over one another.

This cross-over behavior is not evident in stress-strain curves for

stable stainless steels, such as AISI 310. Stable stainless steels

remain completely austenitic in structure, and the elastic moduli and

flow strengths at a given strain progressively increase as test temperature is lowered. In contrast, the metastable AISI 304L steel

stress-strain curves exhibit early departures from linear elasticity at

100 K (and also at 76 K) , so that the proportional limit is lower at

135
these temperatures than at 199 K. In a study of Fe-18Cr-10Ni stainless

steel, an observed decrease of the proportional limit between 243 and

76 K was attributed to e martensite transformation [1]. In these

steels, e forms first, followed by a

1 . Vet strain-induced a1 phase

transformation was confirmed in the present tests of AISI 304L (see

Table 1), and Olsen and Azrin [2] have shown a strong correlation between the stress to initiate a

1 and a . Furthermore, the higher strainhardening rates at larger strains have always been correlated with a'

formation; in this range of the stress-strain curve the volume fraction

of a' is linearly related to strain. Hence at strains above 0.011 for

this alloy, there is the gradual transition to a monotonic temperature

dependence of flow strengths. Clearly, there is currently conflicting

evidence regarding the roles of e and a

1 on flow strength. The evidence

suggests a dual role or possible dual product form for a': at low

strains a' forms at slipband intersections and acts as a "window" for

cross-slip, and at higher strains the a' forms within slipbands and acts

as a dislocation or deformation barrier.

Creep

Application of a constant stress at constant temperature (76 K)

resulted in isothermal martensitic transformation and elongation of

AISI 304L specimens. It has been reported that creep occurs in AISI 300

series steels at temperatures as low as 20 K [5]. The purpose of our

creep tests was to obtain observations at 295 and 76 K to aid interpretations of the thermal cycling tests described later.

Figure 3 shows a series of creep tests performed at three temperatures. No creep occurred at 4 K at the stress levels indicated, but

136
significant creep occurred at 76 and 295 K. ^e creep at 76 K was accompanied by a

1 martensite formation, a indicated in Table 1.

Thermal cycling at constant stress

One AISI 304L specimen was thermally cycled 10 times betwepn 295

and 76 K with no externally applied stress, but no measurable amount of

a' formed. Two additional specimens were "hermally cycled 10 times

while constant stresses ot 100 or 200 MPa wer >. applied A stress of

100 MPa falls in the elastic region, whereas a stress of 200 MPa causes

a small plastic strain of 0.0005. Measurable a' martensite formed in

both specimens and there was also a significant elongation of the specimens. As shown in Table 1, greater elongation and martensitic transformation occurred in the specimen at 200 MPa, but the ratio of perxent

a' martensite ic elongation wdo o factor of three higher for the

specimen held at 100 MPa.

The thermal cycling test at 200 MPa took 0.5 h to complete. In

comparison, more a' phase transformation and specimen extension resulted

from thermal cycling (295 to 76 K) at constant stress than from the

0.5-h, 76 K creep test at the same stress level. However, the ratio of

a' to specimen strain was reduced when thermal cycling was combined with

creep at 200 MPa (Table 1).

Recovery at room temperature

Figure 4 illustrates a novelty discovered in our tests. The usual

strain-hardening behavior for an austenitic stainless steel is depicted

by data from specimen 1, which was monotonical ly loaded in tension at

76 K. Specimen 2 was (1) cooled to 76 K, (2) strained into the plastic

range, (3) completely unloaded, (4) warmed to 295 K. (5) cooled to 76 K,

137
and (6) reloaded at 76 K. The stress-strai n curve on reloading at 76 K

diverges from the previous unloading line, and plastic flow resumes at a

lower stress than expected. In essence, the strain-hardening of the

previous loading cycle is partially erased by the intermediate thermal

cycle (76 to 295 to 76 K) at zero stress. The proportional limit,

observed on each reloading, remains nearly equivalent with each thermal

cycle and this can be repeated for a number of cycles.

The effect is apparently linked with mi crostructural instability.

When the same experiments were repeated using metastable AISI 316 and

stable AISI 310 steels, the anomalous behavior was observed for metastable AISI 316, but not for stable AISI 310. Therefore, to observe the

effect it is necessary (1) to have a metastable steel and (2) to include

the thermal cycle to 295 K and back to low temperature.

After a total elongation of 5%, specimen 1 (monotonic loading)

contained 5% a' martensite, whereas specimen 2 (repeated stress and

thermal cycling) contained 6% a' martensite. This tends to confirm the

suspicion that a

1 martensite forms during each loading cycle at about

the same stress level and that warming to room temperature provides

sufficient internal recovery to permit additional a

1 transformation.

However, this explanation does complicate the earlier proposed concept

that early a' transformation occurs at cross-slip sites, where sufficient dislocation pile-up stresses are presumably generated to induce

transformation . Linder such circumstances, recovery would not be expected to encourage transformation.

SUMMARY

The results of experiments on metastable AISI 304L stainless steels

are

:

138
1. Significant creep occurs at 295 and 76 K, but not at 4 K, at

the stress ranges studied.

2. Thermal cycling between 295 and 76 K serves to increase the

creep rate.

3. Warming to room temperature produces limited recovery after

tensile deformation at 76 K.

4. The role of a

1 and e martensite transformations in the deformation processes remains unclear.

ACKNOWLEDGMENT

The assistance of D. Baxter of NBS in performing tensile tests is

gratefully acknowledged.

REFERENCES

1. Suzuki, T., Kojima, H., Suzuki, K., Hashimoto, T. } Koike, S., and

Ichihara, M., Plastic deformation and martensitic transformation in

an iron-base alloy, Scr. Metal 1. 1_0 (1976), pp. 353-358.

2. Olsen, G. B. and Azrin, M., Transformation behavior of TRIP steels.

Metal! . Trans. A , 9A (1 978) 713-721.

3. Read, D. T., Reed, R. P., and Schramm, R. E. (1979), Low temperature deformation of Fe-18Cr-8Ni steel, Metall. Trans, (to be pub1 ished)

.

4. Reed, R. P. and Mikesell, R. P., The stability of austenitic

stainless steels at low temperatures as determined by magnetic

measurements, in Advances in Cryogenic Engineering , Vol . 4, K. D.

Timmerhaus, ed. (Plenum, New York, 1960), pp. 84-100.

139
Voyer, R. and Weil, L., Creep properties of 18-10 and 25-20 stainless steels at cryogenic temperatures, in: Advances in Cryogenic

Engi neeri ng , Vol . 10, K. D. Timmerhaus, ed. (Plenum, New York,

1 965), pp. 110-116.

140
Table

1.

Summary

of

test

results.

03 4-> O

03

+-> O co LO CXJ

•1

CJ CO rx LO

4-> o . o o . o • • 03 t— o o cnc: oI

LUco

•r— lo d- 1— o d- CD

4-> ^ CXl o CXJ oo 1— o O CXI 03 ^ . • • • • • • cn— o o o o o o o o co

f—

LlJ

oj

4-J • i

U~1 CO o CD o IX LO C c-x o o o o o o o i—

aj • • . • • • . * +j — o o o o o o o o

s_id s:

aj LO LO LO LO E <=C =c <c - " <c • • °r- x: 2: o o ~'Z o o 1— x_-

co *

CO '—

O) rd o o lx o o o o

S- Q- C\J o CO o o o o o

-m s: CNJ CXI CXI CXI cxj 1

cxj co

aj

S_o

-m co co CO

id -•

—•

ir> co co LO CO rx IX IX

S_ iX CTl IX rx CD IX x^ x^ x^

aj -

CXJ cxi LO LO LO

Cl. CD OD cn LO E CXJ CXJ CXI CD

<D CO

1— . co

>> X2

aj • TJ O • 1

>• 1 • X CO 

CXJ d" CXJ CO

d- 13 o 1

1

1—

aj Cl co- •1—

co o o D co

L*C

4-> -M -i-> c c sc o

(d id rd 4-> 4-> +-> +-> CO CO CO id C C £= CL o o O

CD CD CD

a CJ CJ c c: C E •1

•i

— •r— * n *1 o

-a "O "O "O CO CO CO Ca) id id ra aj aj aj 4- E o o o r— i

i

S- 1— i— I

O o a +J o >> >1 >1 >i >> S- 4- cj a <_> 1

r—• a CJ o QJ

S- •1

i

•r— c c >

aj c c £C o o n— 1r—

C

Q- o o o ra id id o +-> +-> +-> CL CL E E E CJ

4-> o o o a> O) S- S_ S_ CO co c c aj QJ aj QJ aj o

QJ o o o S- s- x: sz x: 1

1— se: o c_> t— \— 1

*

141
LIST OF FIGURES

Figure I. Yield strength of AISI 304L stainless steels, showing irregular temperature dependence.

Figure 2. Stress-strain curves for AISI 304L steel, showing cross-over

behavior

.

Figure 3. Creep test results for AISI 304L at three temperatures.

Figure 4. Stress-strain curves for AISI 304L specimens illustrating

recovery effect after unloading and warming to 295 K.

142
YIELD

STRENGTH,

OT

(MPa)

AISI 304 L STAINLESS STEELS

Figure 1. Yield strength of AISI 304L stainless steels, showing irregular temperature dependence.

143
AISI 304L

Fe -18.5 Cr -8.75 Ni -0.03C-0.06IN

Figure 2. Stress-strain curves for AISI 304L steel, showing cross-over

behavi or

.

144
ELONGATION,

%

CREEP TEST

Fe - 18.3 Cr - IO.I Ni -0.028C-0.039N

Figure 3. Creep test results for AISI 304L at three temperatures

.

145
AISI 304 L

Fe -l8.3Cr-IO.INi- 0.028C -0.039N

Figure 4. Stress-strai n curves for AISI 304L specimens illustrating re- covery effect after unloading and warming to 295 K.

146
LOW-TEMPERATURE DEPENDENCE OF YIELDING IN AISI 316 STAINLESS STEELS

R. L. Tobler and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Tensile tests at temperatures between 323 and 4 K were performed on

one heat of AISI 316 austenitic stainless steel having the composition

Fe-17.34Cr-12.17Ni-l.55Mn-2.16Mo-0.051C. The temperature dependences of

the yield and flow strengths at plastic strain increments from 0.2 to

3.65% are analyzed- At the yield strain (0.2%), no body-centered cubic

(bcc) martensite phase transformation is detected. At higher strains

(~3.2 ± 0.6%), bcc martensite forms from the parent austenite phase at

test temperatures below 190 K, but there are no discontinuities in the

temperature dependence of flow strength. A review of data available for

three heats of AISI 316 at temperatures between 973 and 4 K reveals that

deviations from thermally activated plastic flow theory occur at temperatures below 175 K, apparently depending on heat-to-heat compositional

variations. Grain size and magnetic transition effects on the yield

strength are discussed.

147

INTRODUCTION

The mechanical behavior of austenitic stainless steels over a broad

range of temperatures is a subject of interest to designers of superconducting and nuclear energy devices. In this paper, the initial

yielding and stress-strain behavior of AISI 316 stainless steel are

described. AISI 316 stainless steel is a metastable Fe, 16-1 8Cr , 10-

1 4N i steel that may undergo a partial austeni te-to-martensi te phase

transformation at low temperatures, the probability of which increases

with increasing strain, decreasing temperature, and decreasing alloy

content.

Previous studies of Fe-Cr-Ni austenitic stainless steels show that

two martensitic phases may be formed: the hexagonal close packed (hep)

e martensite and the body-centered cubic (bcc) a' martensite. The e

phase usually forms first; then a' appears [1], The a

1 phase is ferromagnetic and easy to detect using a magnetometer [2]. For AISI 316, it

is reported that no a' forms spontaneously on cooling to 4 K, but up to

50% a' forms on straining to failure at 4 K [3].

There is apparently conflicting evidence on the role of martensitic

phase transformation in affecting the low temperature flow behavior of

metastable stainless steels. Suzuki et al.,[4] measuring the temperature dependence of the 0.2% offset yield strength in an Fe-18Cr-9Ni

stainless steel, report a significant decrease in flow strength when the

test temperature is lowered from about 260 K. They argue that e

martensite forms first, following an inverse stress dependence at lower

temperatures, which results in reduced flow strength. Following e martensite formation, a' martensite forms at slip plane intersections and

149
the a' lathes along <110> act as windows for dislocation pileup release;

the result is an "easy glide" region of the stress-strain curve. This

"easy glide" region had previously been reported by Reed and Guntner [5],

who argued that it was caused by strain-induced e martensite, since

their x-ray and transmission electron microscopy data for an Fe-18Cr8Ni alloy indicated an increasing concentration of hep during easy

glide. However, Olson and Azrin [6], studying an Fe-9Cr-8Ni alloy,

found an excellent correlation between flow strength reduction at lower

temperatures and the initiation of a

1 martensite. No "easy glide"

region of the stress-strain curve was observed.

The effect of grain size on low-temperature flow strength of austenitic stainless steels has not been documented. The characterization

of grain size effects is significant in the design of strong and tough

steels for low-temperature applications.

Soviet authors [7,8] report the existence of three major anomalies

in the temperature dependence of the yield strength for Fe-18Cr-xNi

austenitic stainless steels containing 8, 10, or 20% Ni. These anomalies

amount to sharp decreases of the yield strength on the order of 100 MPa.

The first anomaly occurs between 20 and 4 K and is unexplained. The

second anomaly occurs at temperatures between 60 and 35 K and is attributed to the Neel transition (magnetic ordering). The third anomaly

occurs in metastable alloys (lower Ni) at temperatures just above the M<-

temperature and is attributed to strain-induced martensitic phase transformations. Clearly, continuous flow-strength measurements as a function of temperature should be made to assess these results. Therefore,

tensile tests were conducted on a grade of AISI 316 austenitic stainless

steel to assess the influence of martensitic transformation , grain size,

magnetic transitions, and very low temperature gh flow strength.

150
MATERIALS

The AISI 316 stainless steel tested in this study was purchased in

the form of a 38-mm-thick plate. Tensile specimens were machined in the

longitudinal orientation in the form of flat coupons having a 50.8-mm

gage length and a 2.54- x 4.77-mm-gage cross section. After the specimens

were annealed at 1338 K for 0.5 h, they had a Rockwell B (Rg) hardness

of 71 and an average grain diameter of 55 ym (ASTM grain size number

5.5). The chemical compositions of this steel (denoted heat 2) and

other AISI 316 steels referred in this paper [9-11] are listed in Table 1.

PROCEDURE

Tensile tests were performed on (1) annealed and (2) strained and

recrystall ized material. The strained and recrystal 1 ized material was

used to determine grain size effects. All tensile tests were performed

at a crosshead rate of 8.5 ym/s using a screw-driven machine and cryostat

[1,12]. Test temperatures of 4 and 76 K were achieved with baths of

liquid helium and liquid nitrogen. Other temperatures were achieved by

controlled thermal conduction using liquid nitrogen reservoirs and

heating wires attached to the specimen grips and load train. Current to

the heating wires was varied until thermal equilibrium was reached at

the specified test temperature, which was accurate to ± 1.5 K. Loadversus-defl ection curves were recorded with a commercial load cell and

clip-gage extensometer

.

Clip-gage extensometers magnify the specimen strain and permit

measurement of only the specimen strain within the (2,5-cm) gage length.

Estimated flow strength measurement inaccuracies are ± 2%, but lowtemperature specimen-to-specimen variations (including both material and

151
measurement variations) may reach ± 7.5%. After plastic strains of 0.2

to 3.65% the specimens were unloaded, warmed to 295 K, and examined for

the presence of ferromagnetic a' martensite with a torsion-bar magnetometer [2]. X-ray diffractometer measurements were used to try to

detect small concentrations of e martensite. No e martensite was detected and our estimated sensitivity is 2%. The stress-strain curves

were calculated in the conventional manner, and the measurement uncertainty of flow strength values is not greater than 3%. As an indication of data scatter, seven tests were performed on the AISI 316 steel

at 76 K, resulting in an average yield strength (a ) of 448 MPa and an

overall spread of ± 7 MPa.

Grain size variations were obtained by recrystal 1 i zation treatments. Tensile specimens were plastically strained 20 to 30% at 295 or

76 K and then annealed for 0.25 to 0.5 h in an argon atmosphere at 1233

to 1573 K. These treatments produced grain sizes ranging from 22 to

140 ym. Subsequent to recrystal 1 ization treatments, the specimen surfaces were cleaned by electropolishing. The recrystal 1 i zation treatments

increased the scatter of yield strength measurements to about ± 15 MPa.

TENSILE YIELD STRENGTH RESULTS

The tensile yield strength (a ) and initial stress-strain behavior

of the Fe-17Cr-12Ni-2Mo alloy were measured as a function of temperature, grain size, martensitic transformations, and magnetic transitions.

These effects are described in this section.

Temperature : As listed in Table 2, twenty-six a measurements were

obtained for annealed AISI 316 stainless steel (heat 2) at seventeen

temperatures between 323 and 4 K. Magnetometer inspections of the
specimens indicated no a

1 martensitic transformation for this heat after

0.2% plastic strain, regardless of test temperature. The a data between

295 and 4 K are plotted and compared with previously published data [9,10,

13-16] in Fig. 1. This comparison reveals significant heat-to-heat

variations of o^, with considerably greater differences at cryogenic

temperatures than at room temperature.

Several references cited in Fig. 1 report data at only a few test

temperatures, whereas the functional dependence of a on temperature can

be identified only by extensive measurements performed over a broad

temperature range. Therefore, data for three heats of AISI 316 are

replotted in Fig. 2. These data show that a increases progressi vely as

temperature is lowered. Despite minor variations of grain size, strain

rate, and composition, the a values for the three heats are nearly

equivalent between 973 and 175 K, where the test temperature intervals

from different studies overlap. Below 175 K, significant variations of

strength begin to emerge. At still lower temperatures , the variations

become more obvious until, at 4 K, correlates with composition: Fe17. 250-13. 48Ni (heat 1) exhibits the highest strength, Fe-17.34Cr1 2 . 1 7N i (heat 2) exhibits intermediate strength, and Fe-1 6. 72 Cr- 1 1 . 2 7N

i

(heat 3) exhibits the lowest strength. This apparent association of

higher strength with higher nickel content may relate to crystal structure

stability as discussed below.

Thermally activated plastic flow theory predicts that should

exhibit some dependence upon strain rate. A few experiments were performed on the annealed AISI 316 stainless steel to evaluate this effect.

Increasing the crosshead rate by a factor of 100 from 0.85 ym/s to

85 ym/s increased the flow strength at strains lower than 1% by

153
factors of 6 % at 295 K and 3 % at 76 K. The effect of crosshead rate on

the a at 4 K was negligible. Therefore, strain rate effects do not

play a significant part in explaining the variability noted in Fig. 1.

Strength variations may arise between specimens from a single heat

of steel owing to inhomogeneity or thermomechanical processing effects.

As shown in Fig. 1, Read and Reed's data [13] for a 316 stainless steel

at 295, 76, and 4 K are 5 to 17% higher compared with the results of

Tobler, Reed, and Burkhalter [9], who tested the same heat. In the

former study, the specimens were tested in the as-machined condition,

whereas the specimens used in the second study were annealed (1289 K,

1 h) after machining. This annealing step reduced the hardness at 295 K

from Rg79 to R^7 5 and lowered a by increasing the grain size slightly

and removing any residual cold work from the fabrication or machining

processes

.

An example of within-heat variability is also noted in the present

study: the annealed AISI 316 steel specimens had a 76-K a of 448 ±

7 MPa, whereas reprocessed (strained and recrystallized) specimens

having a similar grain size (55 ym) had a 76-K a of about 480 ± 15 MPa.

Thus, process history influences strength. Larger scatter for the

recrystal 1 ized specimens is to be expected since dislocation density and

substructure within the grains depend on prior history and recrystallization temperature.

Grain Size : Figure 3 illustrates the effect of grain size on a at

295, 76, and 4 K. The data at each temperature fit an inverse square

root dependence on the average grain diameter, d, which is in reasonable

accord with the familiar Hall-Petch equation:

154
a = a . + kd

y l

( 1 )

- 1/2

where a. refers to the flow strength in the absence of grain boundaries

- 1/2

(the "friction stress") and k is the slope of the a^-versus-d plot

and indicates the grain boundary dependence of the yield strength. The

values of and k are given in the figure.

Martensitic transformation : Specimens from this study were screened

using a magnetometer, but no a

1 martensite was detected after 0.2%

plastic strain. To force the phase transformation to occur, a series of

specimens was loaded to higher strains. Figure 4 presents the strain

dependence of a' at 76 K. Notice that a plastic strain of about 2% is

necessary for a' to form at 76 K. The dependence of a' on plastic

deformation is linear, and this trend continues at least to a strain of

5%. At larger strains, the dependence apparently increases since fractured

specimens show between 44 and 52% a' at uniform strains of 40 to 50%.

The stress-strain curves at higher strain increments are presented

in Figs. 5 and 6. The quantity of a' martensite formed after either 3.65

or 2.75% plastic strain is negligible for test temperatures between 300

and 190 K, but below 190 K, a' is formed in increasing amounts. In

fact, as illustrated in Fig. 7, the percent a

1 martensite divided by the

strain is linearly dependent on test temperature. The curves of Fig. 5

and 6 represent specimens in various states of transformation, containing

from 0 to 0.1% a' martensite per unit strain. But the stress-strain

curves are regular and free from discontinuities, indicating that the

martensitic transformation occurs gradually with no anomalous effects.

These findings contrast with the behavior of Fe-18Cr-8Ni and Fe-18Cr1 ON i stainless steels that exhibit: (1) irregular yield and flow strength

155
temperature dependences owing to the initiation of e and a

1 phase

transformations [7], and (2) suppression of a

1 martensite at 4 and 20 K,

compared with that at 76 K [5].

DISCUSSION

Thermally activated flow : According to classical theory (e.g.,

Seeger [17]), a may be separated into two components:

CT y

+ o- ( 2 )

Here a p

is an athermal stress component determined by metallurgical

structure, and a-j- is a thermal stress component governed by lattice

vibrations and rate processes. By estimating from Fig. 2, the value of

Op (uncorrected for shear modulus) for the AISI 316 steel heats 1 and 2

is taken to be 100 MPa, which is the asymptote approached by a at

elevated temperatures. From previous work involving thermodynamic and

thermally activated plastic flow theories, the following expression was

developed for Oji

a y

= (a

Q

- Op) exp (- CT) (3)

Here Oq is the extrapolated value of the yield strength at absolute zero

and C is a materials constant comprised of activation energy, entropy,

and strain-rate terms. This expression derived from the work of Yaroshevich

and Ryvkina, who did not correct the Op value for the effect of shear

modulus [18].

The applicability of Eq. (3) for AISI 316 stainless steels was

tested by plotting log (a - op) versus T. Data agreeing with Eq. (3)

will produce a single straight line when graphed this way, assuming that

156
a single mechanism of plastic flow is operative. As demonstrated in

Pig. 8, Eq. (3) successfully predicts the experimental results for the

three AISI 316 steels down to 175 K, but divergences from the main trend

occur for each of the three heats at 175, 150, or 30 K, depending on

alloy composition. A graphical solution leads to the following expression for the AISI 316 data (heats 1 and 2) between 975 K and the

temperature of divergence:

a y

= 100 + 700 exp (-6.2 x 10“ 3

T) (4)

As a consequence of divergences from the main trend, the experimentally

measured 4-K a values for all heats are lower than predicted by Eq. (4).

If the trend lines of Fig. 8 are extrapolated, then the measured

values for heats 1 and 2 at 4 K (near absolute zero) are 15 to 45% lower

than the extrapolated value of ag, which is 780 MPa. The data for

heat 3 suggest an even greater difference.

The reasons for deviations from Eq. (4) are not completely understood. It was suggested previously that dislocation tunneling and

magnetic ordering may affect a at temperatures approaching absolute

zero [9]. Likewise, adiabatic heating becomes significant as absolute

zero is approached. But these processes alone are unable to account for

the deviant behavior of AISI 316 heats 2 and 3 at temperatures in the

150- to 175-K range. Moreover, these processes do not account for the

apparent sensitivity to alloy composition. As noted in Fig. 1, divergences from Eq. (4) for heats 1, 2, and 3 occur at 30, 150, and 175 K,

respectively. This sequence follows the order of decreasing alloy

content (Table 1), decreasing stacking fault energy (Table 3), and

157
decreasing crystal structure stability, as indicated by the and M

Q

temperatures (Table 3). The possible martensite roles are discussed in

the next section. It is also possible that more than one dislocation

mechanism is operative, depending on temperature and composition. If a

second thermally activated dislocation process becomes rate controlling

at sufficiently low temperatures, the change of slopes in Fig. 5 would

result.

In Figure 9, the data for the three AISI 316 stainless steels

have been normalized by using the shear modulus and plotted versus test

temperature, which is normalized by using the melting point, taken as

T^ = 1658 K. Values of the shear modulus at cryogenic temperature were

taken from Ledbetter [11]. These data indicate that the thermal component

of a

y

becomes active at a temperature of 1/3 the melting point. A

linear extrapolation of the thermal portion of the yield strength gives

a value of about 300 K for the intersection of the two components. The

divergences from the thermal components at low temperatures , shown in

Figure 8, is also observed here.

Role of martensitic transformation : The nonlinear dependence of

In a-j. (Fig. 8) or on (Fig. 9) on temperature at very low temperatures

may result from strain- induced martensitic transformation

.

The and temperatures in Table 3 provide only a relative

indication of stability, having been calculated on the basis of the Cr,

N i , Mn, Si, C, and Mo contents. The calculated M<~ values are negative,

indicating complete stability on cooling to absolute zero, which is

consistent with the observation that no a' martensite formed on cooling

158
to 4 K. Beyond this, a quantitative comparison of the measured and

calculated transformation temperatures is impractical because the calculated and temperatures are influenced by trace elements which

are neither controlled nor analyzed during steel manufacture.

It is unlikely that the a

1 transformation can account for the

observed deviations since no a' martensite was detected at 0.2% plastic

strain. Furthermore, from Fig. 4 it is apparent that no a' is expected

to form until a plastic strain of about 2% is reached. However, it is

possible that martensitic transformations producing small amounts of e

phase may be responsible. From Table 3, the calculated stacking fault

energies decrease, but remain positive, with increased deviation of a

from linearity. This possibility could not be confirmed, since our

metal lographic and x-ray inspection methods lacked the required sensitivity to detect a small amount (< 2 %) of e.

In the commercial Fe-18Cr-8Ni austenitic stainless steel grade

(AISI 304), the a

1 martensitic product and, perhaps, the e martensitic

product apparently contribute to the anomalous temperature dependence of

0 ^ and to the "easy glide" portion of the stress-strain curves at temperatures near 77 K [4,5,23]. However, the data represented in Figs. 4-7

suggest that in the alloy of this study (Fe-1 7 Cr- 1 2Ni-2Mo) the temperature dependence of the flow strength (including a ) is well behaved and

the a' martensitic product, therefore, does not contribute to anomalous

behavior

.

The nature of the linear relationship between the "normalized"

percent bcc martensite (percent magnetic martensite divided by engineering plastic strain) and temperature deserves additional experimental
clarification. It is implied that the temperature of about 190 K represents the onset of strain-induced martensite, and at higher strains and

temperatures little magnetic product phase is detected. This behavior

has been substantiated by the tests to fracture at 195 and 295 K. However, the nature of the probable family of curves, generated by taking

data at larger strains, is not clear. Similar data for AISI 304 steel

indicate a gradual deviation from linearity at larger strains, but the

existence of a constant temperature of a' initiation [24].

With respect to the role of martensitic products in affecting the

temperature dependence of the flow strength of Fe-Cr-Ni alloys at low

temperatures, there are two schools of thought: Suzuki et al . [4] argue

that the e phase must be responsible for the reduction of a at low

temperatures since in their study a

1 martensite was detected at strengths

considerably greater than a . Additionally, they have argued from

optical metal 1 ography evidence that e martensite forms at strength

levels near the measured a . Conversely, Olson and Azrin [6] studied a

less stable Fe-9Cr-8Ni alloy and obtained excellent correlation between

the temperature dependence of o and the flow strength at which a' martensite was first detected. From the results of this study, it is

apparent that the onset of a' strain-induced martensite in austenitic

Fe-Cr-Ni alloys does not lead to a reduction of the flow strength at low

temperatures. This finding, again, raises the question of the nature of

the a', ferromagnetic martensitic product or products that form(s)

during plastic deformation. Conflicting parametric studies and lack of

conclusive microscopy preclude assessment of the role of e and a' martensites in affecting low temperature flow strength.

160
Role of grain size : Data on the effect of grain size for austenitic stainless steels are rare. The present data at room temperature

indicate a very slight effect, as evidenced by the low-k value of

- 1/2 - 1/2

191 MPa*ym . This value is in agreement with the 221 MPa*ym

value obtained by Norstrom [25] for a 31 6L stainless steel containing no

nitrogen. Low-k values are typical for face-centered cubic alloys at

room temperature. At cryogenic temperatures, however, the grain size

effect increases significantly. As shown in Fig. 3, the values of a.

and k at 4 K are greater by factors of 2 and 5, respectively, than the

values at 295 K.

In Fig. 10, the temperature dependences of k and a. from Eq. (1)

are plotted, using the 4- to 295- K data from this study and the 295- to

875-K data of Norstrom [25]. The k is strongly temperature-dependent at

low temperatures, but insensitive to temperature above room temperature.

The frictional stress, a., increases with decreasing temperature and the

In a. is linearly dependent on T(K) for temperatures lower than about

500 K. These data clearly imply that grain size has a significant

effect on low- temperature flow strength. From Fig. 10, the quantity ay

-

a. represents the grain boundary contribution to flow strength and about

30% of 0 ^at 4 K is contributed by grai n-boundary strengthening.

Role of magnetic transition : Similar to other Fe-Cr-Ni austenitic

stainless steels, AISI 316 undergoes a magnetic ordering phenomenon

characterized by a change from the paramagnetic to the antiferromagnetic

state at low temperatures. Accordingly, a peak occurs in the temperature dependence of magnetic susceptibility at T^, the Neel transition

temperature, which corresponds to the onset of antiferromagnetic ordering. Since a decrease of elastic constants occurs at T^ [26,27], it

161
was suggested that the Neel transition likewise affects dislocation

motion and hence tensile properties [7,8]. On the basis of Soviet

experimental data [7,8], a sharp decrease of of about 100 MPa should

be expected for our Fe-18Cr-12Ni alloy at temperatures between 60 and

48 K. As shown in Fig. 11, elastic property measurements for Fe-17.5Cr1 2 . 9N i stainless steel do reveal a noticeable drop in Young's modulus

(E), identifying the Neel transition between 40 and 20 K. From Table 3,

the calculated predicted value of is 26 K, but the data of Fig. 11

indicate no decrease of a at temperatures corresponding to T^. From

Table 3, T^ values for the three AISI 316 steels of Fig. 8 are nearly

equivalent, despite minor compositional differences.

Care must be exercised in these arguments, however. The approximate reduction of E after magnetic transition is about 2%, which is

within the possible a data variation. For better assessment, the large

grain y size contribution should be subtracted from a , and a. be examined

y i carefully as a function of temperature.

SUMMARY

Tensile tests were performed on AISI 316 stainless steel (Fe17.34Cr-12.17Ni ) at many temperatures between 323 and 4 K, and the

results were compared with data for other AISI 316 heats at temperatures

to 973 K. The findings are:

1) The existing data for AISI 316 show a monotonic temperature

dependence, with a ^ increasing continuously as temperature is

lowered between 973 and 4 K.

162
2) Between 973 and 175 K, data for AISI 316 follow a single

equation derived from the theory of thermally activated plastic

flow. Below 175 K, deviations occur, apparently depending on

chemistry variations which in turn affect phase stability and

dislocation dynamics.

3) Deformation beyond 0.2% plastic strain is needed to initiate

transformation to a

1 martensite at temperatures below 180 K.

4) Reduction of grain size contributes significantly (30% at 4 K)

to low-temperature flow strength and at low temperatures the grain

size dependence of a^

follows a Hall-Petch relationship.

5) Magnetic transitions do not noticeably affect the deformation

of AISI 316.

ACKNOWLEDGMENTS

We thank D. Beard, Office of Fusion Energy, DoE for his support and

continued encouragement. D. Beekman and D. Burkhalter of NBS contributed

significantly to this study by conducting most of the tensile tests. Discussions with Dr. E. L. Brown of NBS, who contributed valuable metal 1 ographic

assistance and reviews of the manuscript, are also gratefully acknowledged.

REFERENCES

1. Read, D. T. , Reed, R. P. , and Schramm, R. E., Low temperature

deformation of Fe-18Cr-8Ni steels, in: Materials Studies for

Magnetic Fusion Energy Applications at Low Temperatures-I

I

, NBSIR 79-

1609, F. R. Fickett and R. P. Reed, eds.. National Bureau of Standards,

Boulder, CO (June 1979), pp. 151-172.

2. Reed, R. P. and Mikesell, R. P., The stability of austenitic stainless steels at low temperatures as determined by magnetic measurements, in: Adv. Cryo. Eng. , 4-, K. D. Timmerhaus, ed. Plenum Press,

New York (1960), pp. 84-100.

163
3 . Larbalestier, D. C. and King, H. W., Austenitic stainless steels at

cryogenic temperatures I, structural stability and magnetic properties, Cryogenics , 1 3 (1 973) 1 60-1 68.

4. Suzuki, T., Kojima, H., Suzuki, K., Hashimoto, T, and Ichichara, M.,

An experimental study of the martensite nucleation and growth in

18/8 stainless steel, Acta Metal 1 . , 23 (1972).

5. Reed, R. P. and Guntner, C. J., Stress-induced martensitic transformations in 18Cr-8Ni steel, Trans. AIME , 230 (1969) 1713-1720.

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

Metal! . Trans. A , 9A (1978) 713-721.

7. Ilyichev, V. Ya., Medvedev, Ya. M., Shapovaliev, I. A., and Klimenko,

I. N., Low temperature anomaly of the temperature dependence at the

flow stresses in iron-chromium-nickel alloys, Phys . Met. Metal! ogr. ,

44, No. 1 (1978) 173-176.

8. Verkin, B. I., Ilyichev, V. Ya., and Klimenko, I. N., The lowtemperature change of the magnetic structure and plastic properties

of Fe-Cr-Ni alloys, in: Adv . Cryo. Eng. , 26 Plenum Press, New York

(1980), pp. 120-125.

9. Tobler, R. L., Reed, R. P., and Burkhalter, D. S., Temperature

dependence of yielding in austenitic stainless steels, in: Adv .

Cryo. Eng. , 26 , Plenum Press, New York (1980), pp. 107-113.

10. Sanderson, G. P. and Llewellyn, D. T., Mechanical properties of

standard austenitic stainless steels in the temperature range -196

to +800C, J . Iron Steel Inst

.

, 207 (1 969) 1 129-1140.

11. Ledbetter, H. M., Sound velocities and elastic constants of austenitic stainless steels, to be published.

164
12. Reed, R. P., A cryostat for tensile tests in the temperature range

300 to 4 K, in: Adv . Cryo. Eng. , 7_, K. D. Timmerhaus, ed . , Plenum

Press, New York (1961), pp. 448-454.

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

selected austenitic stainless steels at cryogenic temperatures, to

be published, Cryogenics (1980).

14. Baughman, R. A., Gas atmosphere effects on materials. Progress

Report #2, AF 33(616), General Electric Co. (November 1958).

15. Tobler R. L., Mikesell, R. P., Durcholz, R. L., Fowlkes, C. W., and

Reed, R. P., Fatigue and fracture toughness testing at cryogenic

temperatures, NBSIR 74-359, National Bureau of Standards, Boulder, CO

(March 1974), pp. 182-308.

16. Hoke, J. H., Mabus, P. G., and Goller, G. N., Mechanical properties

of stainless steels at subzero temperatures. Met. Prog. (May 1949)

643-648.

17. Seeger, A., The generation of lattice defects by moving dislocations, and its application to the temperature dependence of the

flow-stress of FCC crystals, Phil. Mag., 4(5 (1 955) 11 94-1 21 7.

18. Yaroshevich, V. D. and Ryvkina, D. G., Thermal -activation nature of

plastic deformation in metals, Sov . Phys.-Solid State (Engl. Transl.)

12, No. 2 (1970) 363-370.

19. Schramm, R. E. and Reed, R. P., Stacking fault energies of seven

commercial austenitic stainless steels. Metal 1 . Trans

.

, 6A (1 975)

1345-1351 .

165
20. Larbal estier , D. C. and King, H. W., Prediction of the low temperature stability of type 304 stainless steel from a room temperature

deformation test, in: Proceedings of the Fourth International Cryo -

genic Engineering Conference , K. Mendelssohn, ed . , IPC Science and

Technology Press, Guildford, Surrey, England (1975) pp. 338-340.

21. Williams, I., Williams, R. G., and Capellaro, R. C., Stability of

austenitic stainless steels between 4 K and 373 K, in: Proceedi ngs

of the Sixth International Cryogenic Engineering Conference , ed

.

K. Mendelssohn, IPC Science and Technology Press, Guildford, Surrey,

England (1976), pp. 337-341.

22. Warnes, L. A. and King, H. W., The low temperature magnetic properties of austenitic Fe-Cr-Ni alloys 2, the prediction of Neel temperatures and maximum susceptibilities. Cryogenics , 1

6

, No. 11

(1976) 659-667.

23. Guntner, C. J. and Reed, R. P., The effect of experimental variables including the martensitic transformation on the low temperature mechanical properties of austenitic stainless steels, Trans

.

Am. Soc . Met

.

, 55 (1 962) 399-41 9.

24. Starr, C. D., Notes on the plastic critical temperature in straininduced martensite reaction, Trans. AIME , 1 97 (1953) 654.

25. Norstrom, L. A., The influence of nitrogen and grain size on yield

strength in Type 316L austenitic stainless steel, Met , Sc i

.

, 1

1

(1977) 208-212.

26. Collings, E. W. and Ledbetter, H. M., Sound velocity anomalies near

the spin glass transition in an austenitic stainless steel alloy,

Phys. Lett. , 72A , No. 1, 11 (1979) 53-56.

1 66
27. Ledbetter, H. M. and Collings, E. W., Low-temperature magnetic

elastic-constant anomalies in three manganese stainless steels

The Metal Science of Stainless Steels , E. W. Collings and H. W

King, eds.

and

, in:

167
Table

1.

Chemical

compositions

(weight

percent)

for

AISI

316

steel

heats

referred

to

in

this

study.

coo g 1 i o 1

CTi i i

o g • 1 o I O 1 1

LO 00 o o CO i

LO oo

2: • • • • OJ OJ OJ CM

co

=5 i • 1 1 o o 1 1

00 CO OJ CD * <

lo LO CO lo oo • • • . o o o o CTi o t

1— C\J o

O o O • • e « o o o

LO o C\J OJ CO o o o

Q_ • O i • o o i o

C\J LO LO LO CD

O o o o o • • • * o o o o

lo LO OJ c CO LO LO • • • 1—

CO hs CO

•i— iOJ cr> • • • • CO OJ 1

C\J r—

LO OJ CD

S- OJ CO "d- o • e • • LO 1^. * 

” *"

<D

CJ

cu

C

— 3 r U- ro t—

<T3 _Q

>> “O

o3 13 O) +-> 1 1 i—

• <_» i—

00 o r— o c CT> ,— r—

SI cu i i LO 1 1 1 1 s_ •r—

+-> CD Jd

<T3 4- +-> O) O) =c cc r— OJ CO

168
Table 2. Yield stress measurements at 0.2%

plastic strain obtained for AISI 316 (heat 2) of this study.

Test Temperature, T,

K Yield Stress, o , MPa* y

323 185

295 204

295 198

295 198

240 244

215 281

194 298

192 306

190 307

164 349

154 368

143 380

125 399

100 428

76 456

76 451

76 445

76 450

76 441

76 447

76 448

60 456

39 487

23 528

4 537

4 527

*To convert from MPa to ksi, divide by 6.895.

169
Table 3. Calculated stacking fault

temperatures for AISI 316

energies (SFE)

steel heats of

and M<~, M^,

this study.

and T

n

Heat No. & Reference SFE

a

, M

b

"s ' M

C D

» T

^

mJ/M

2 K K K

1 [9] 70.4 -226 190 24.7

2 this study 59.6 -126 232 25.5

3 [10] 57.4 - 47 246 7.4

4 [11] NA NA NA 26.4

a. Calculated from equation (9) of Schramm and Reed [19].

b. M<~ is the temperature at which martensite starts forming spontaneously

on cooling, as calculated from the equation of Larbalestier and

King [20].

c. is the temperature above which no deformation induced martensite

is possible, as calculated from the equation of Williams, Williams,

and Capellaro [21].

d. is the temperature of the onset of anti ferromagnetic ordering,

as calculated from the equation of Warnes and King [22].

170
LIST OF FIGURES

Fig. 1

Fig. 2

Fig. 3

Fig. 4

Fig. 5

Fig. 6

Fig. 7

Yield strength of AISI 316 stainless steels at low temperatures, illustrating heat-to-heat variability. Heat 1 is from

Tobler, Reed, and Burkhalter [9]; heat 3 is from Sanderson and

Llewellyn [10].

Yield strength data for AISI 316 stainless steels in the

elevated-to-cryogenic temperature range, illustrating overall

temperature dependence.

Effect of grain size on the yield strength at three test

temperatures for the Fe-1 7Cr-l 2Ni -2Mo alloy.

Dependence of bcc (a

1 ) martensite on tensile plastic strain.

Martensite was detected by using a bar magnet and torsion

balance beam.

Temperature dependence of the yield and flow stresses for

AISI 316 stainless steels, showing no evidence of anomalies

despite martensitic phase transformation occurring at higher

strain increments and lower test temperatures.

Stress-strain curves for AISI 316 stainless steels at various

test temperatures, showing no evidence of anomalies despite

martensitic phase transformations occurring at higher strain

increments and low temperatures.

Amount of bcc martensite per unit strain in AISI 316 specimens

deformed to plastic strains of 2.75 or 3.65%, indicating clear

evidence of martensitic phase transformation at cryogenic

temperatures. Extrapolated temperature of 0% bcc martensite

is about 180 K.

171
Fig. 8.

Fig. 9.

Fig. 10.

Fig. 11.

Temperature dependence of the thermal stress component, illustrating applicability of thermally activated plastic flow

theory at temperatures between 900 and 175 K.

Normalized yield strength versus normalized temperature plot

for three AISI 316 stainless steel heats, showing variations

in thermal and athermal stress components.

Dependence of the tensile yield strength, a^; the Hall-Petch

parameter, K; and the flow strength, a. , for AISI 316 steel.

High temperature data from Norstrom 25 and a. is adjusted to

0.04 weight percent nitrogen.

Temperature dependence of yield stress and Young's modulus for

AISI 316 stainless steel heats, showing no drop in yield

strength at the Neel transition temperature.

172
800 T T

CO Q.

UJ

>

600 -

O 400

LU

DC

HC0

AISI 316 Steel

(annealed)

•Heat 1

= 200 -

Heat 2

.(this study)

Heat 3

A Read and Reed [l 3]

O Baughman [14]

v Tobler, Mikesell, Durcholz,

Fowlkes, and Reed [l 6]

Hoke, Mabus, and Goller [16]

o Sanderson and Llewellyn [10}

100 200

TEMPERATURE, K

300

Fig. 1. Yield strength of AISI 315 stainless steels at low temperatures, illustrating heat-to-heat variability. Heat 1 is from

Tobler, Reed, and Burkhalter [9]; heat 3 is from Sanderson and

Llewellyn [10].

173
Fig. 2 Yield strength data for AISI 316 stainless steels in the

elevated-to-cryogenic temperature range, illustrating overall

temperature dependence.

174
175
%

bcc

MARTENSITE

TENSILE STRAIN

Fig. 4. Dependence of bcc (a 1 ) martensite on tensile plastic strain.

Martensite was detected by using a bar magnet and torsion

balance beam.

176
1000 T T T

TEMPERATURE, K

Fig. 5. Temperature dependence of the yield and flow stresses for

AISI 316 stainless steels, showing no evidence of anomalies

despite martensitic phase transformation occurring at higher

strain increments and lower test temperatures

.

177
FLOW

STRESS,

cr,

MPa

PLASTIC STRAIN, %

Fig. 6. Stress-strain curves for AISI 316 stainless steels at various

test temperatures , showing no evidence of anomalies despite

martensitic phase transformations occurring at higher strain

increments and low temperatures.

178
%

bcc

MARTENSITE

%

STRAIN

TEST TEMPERATURE, K

Fig. 7. Amount of bcc martensite per unit strain in AISI 316 specimens

deformed to plastic strains of 2.75 or 3.65%, indicating clear

evidence of martensitic phase transformation at cryogenic

temperatures. Extrapolated temperature of 0% bcc martensite

is about 180 K.

179
THERMAL

COMPONENT

OF

YIELD

STRENGTH

Fig. 8 Temperature dependence of the thermal stress component, illustrating applicability of thermally activated plastic flow

theory at temperatures between 900 and 175 K.

180
YIELD

STRENGTH

TEST TEMPERATURE, K T

MELTING TEMPERATURE, K

’ TM

Fig. 9. Normalized yield strength versus normalized temperature plot

for three AISI 316 stainless steel heats, showing variations

in thermal and athermal stress components.

181
STRENGTH,

MPa

Fig. 10. Dependence of the tensile yield strength, ay

; the Hall-Petch

parameter, K; and the flow strength, a., for AISI 316 steel.

High temperature data from Norstrom 25 and o

.

is adjusted to

0.04 weight percent nitrogen.

182

GRAIN

SIZE

COEFFICIENT,

K,

MPa

•

/xm
YIELD

STRENGTH,

a;,.

MPa

Fig. 1

TEMPERATURE, K

. Temperature dependence of yield stress and Young's modulus for

AISI 316 stainless steel heats, showing no drop in yield

strength at the Neel transition temperature.

183

YOUNG'S

MODULUS,

E,

GPa

LOW-TEMPERATUP.E DEFORMATION OF 5083-0 ALUMINUM ALLOY

R. L. Tobler and R. P. Reed

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

This paper describes the temperature dependence of the yield

strength and stress-strain characteristics for 5083-0 aluminum alloy.

New experimental data in the temperature range 4 to 340 K are compared

with previously published data extending to the melting point. Thermally activated plastic deformation apparently governs the behavior

between 4 and 400 K, but the athermal component is not simply a function

of the shear modulus. Aspects of serrated yielding and load-versusdisplacement control tensile tests are discussed.

185

INTRODUCTION

Aluminum alloys have been proposed for structural applications in

superconducting magnets for magnetic fusion energy systems. The most

widely used aluminum alloy in cryogenic service today is 5083-0 aluminum, which is used extensively in thick sections for the storage of

liquefied natural gas. This alloy is strengthened by magnesium in solid

solution and is relatively soft, but tough. This provides the advantage,

however, that strength can be maintained in the welded joints of structures made with this alloy.

The low-temperature properties of relatively pure aluminum have

been reviewed by Reed [1], The purpose of this study was to characterize the plastic deformation of 5083-0 alloy over a broad temperature

range. Although scatter in the studied material prevented determination

of the exact form of the temperature dependence of the yield strength,

several observations are recorded concerning the effects of temperature,

magnesium concentration, and method of testing on the 1 ow-temperature

strength of aluminum alloys. This work provides a basis for future

studies

.

PROCEDURE

Material

A 3.8-cm-thick plate of 5083-0 aluminum was obtained from a commercial manufacturer. The ASTM specification B 209-66 gives the chemical

composition in weight percent as follows: 4. 0-4. 9 Mg, 0.3-1 .0 Mn, 0.05-

0.25 Cr, 0.4 max Si, 0.4 max Fe, 0.10 max Cu, 0.25 max Zn, 0.15 max Ti,

and 0.15 max of other impurities. The microstructure consists of elongated grains in the rolling direction, as illustrated in Figure 1.

Typical Rockwell hardness of the alloy is Rg 61.

187
Specimen

Tensile tests were performed using unnotched, longitudinally

oriented specimens having a 50.8-mm gage length and 2.54 x 4.80 ym

cross section. Specimens were cut from the 3.8-cm-thick plate (at

various thickness depths), then annealed at 686 K and furnace cooled.

Apparatus

Most tests were performed using the screw-driven machine and variable temperature cryostat described by Read and Reed [2]. The crosshead

rate during these tests was 8.5 ym/s. The tensile yield strength was

recorded at 0.002 offset specimen strain. The specified test temperatures ranged from 4 to 340 K and were accurate to better than ± 1 .5 K.

A few tests were also performed using a servo-hydraulic test machine in

load control at a loading rate of 0.1 kN/s. Tests were performed at

temperatures between 4 and 76 K in a gaseous helium environment and

between 76 and 295 K in a gaseous nitrogen atmosphere.

RESULTS AND DISCUSSION

Tensile Yield Strength

In Figure 2, the tensile yield strength between 4 and 300 K is

presented. The data scatter is large, with six measurements of the

room-temperature yield strength ranging from 137 MPa to 158 MPa. It is

thought that this scatter is real and reflects inhomogeneity in the

3.8-cm-thick 5083-0 plate. Unfortunately , the scatter is too large to

permit accurate interpretation of the temperature dependence at low

temperatures . The trend shown in Figure 2 follows the minimum values

measured in this study.

188
The present results are combined with high temperature results [3]

in Figure 3. In the temperature range 150 to 400 K, the yield strength

(a ) is constant, if one relies on the handbook values in this temperature interval. Normalizing the yield strength by dividing by the shear

modulus (G), which is reported by Weston, Naimon, and Ledbetter [4],

results in a negative temperature dependence of a /G in this temperature

range (Fig. 3). The rather strong thermal contribution to the flow

strength is most likely associated with dynamic solute (Mg) segregation

to dislocations. Evidence for this effect is the discontinuous yielding

observed in the stress-versus-strai n behavior of the alloy at room

temperature that disappears at 76 K. Owing to the rather strong dynamic

strain aging contribution, it is not possible from these data to separate the athermal from the thermal components.

Stress-Strain Behavior

True stress-true strain curves for the 5083-0 aluminum alloy at

295, 76, and 4 K are presented in Figure 4. Here true stress is defined

L-L L-L

as P/A

q

(1 +

l~°

) anc* true strain equals In (1 + -

j

-- - ) where P is load,

o o A

q is the original cross-sectional area, L q is the original gage length,

and L is the instantaneous gage length. Serrated yielding phenomena

were observed at 4 K and at 295 K, but not at 76 K. In the case of

5083-0 alloy at 4 K, the stress-strai n curve was smooth up to a strain

of about 4%, at which point the flow became discontinuous. The stressstrain curve at 295 K showed single serrations at 1.8 and 2.5% strain,

then continuous serrations from 2.7% strain to fracture.

The serrated yielding at 295 K is well known in aluminum-magnesium

alloys as the Portevain-LeChatel ier effect, which is accompanied by

189
Luders band propagation and is attributed to dynamic strain aging [5].

In the study reported by Lloyd [6], dynamic strain aging was observed at

temperatures between 200 and 350 K. In our study, the stress-strain

curve at 295 K showed continuous serrations from 3% plastic strain to

fracture. The strain at which the serrations begin and the magnitude of

the load drops associated with each serration have been demonstrated to

be dependent on strain rate, magnesium concentration, and temperature [7].

The serrated yielding at 4 K is attributed to adiabatic heating, a

completely different mechanism from that of dynamic strain aging at

295 K. Adiabatic heating is possible because of the low specific heat

of the alloy at 4 K [8]. Thus the amount of heat necessary to produce

an appreciable temperature rise is relatively low. As a result of

heating during plastic deformation, appreciable softening occurs and the

deformation becomes unstable. Low thermal conductivity at 4 K (compared

with that at 295 K) contributes to a localization of the heat produced

until flow occurs.

During testing at 4 K, the load drops suddenly and the flow becomes

localized when thermal softening, due to heat generated in plastic

straining, is greater than the strain-rate hardening. This instability

criterion is more easily satisfied at the lowest temperatures because

1) the volume specific heat is reduced so that relatively large temperature increases are possible in straining and 2) the flow stress is more

strongly temperature dependent.

Specimens tested using load control at 4 K were compared with those

tested using stroke control. The stress-strain curves obtained are

compared in Figure 5. The adiabatic instabilities corresponded to

specimen strains of the order of 0.04%. The load-controlled tests did

190
not produce catastrophic deformation during the discontinuous yields.

Instead, each local serration is arrested at some strain and temperature

where the flow process becomes stable, owing to local specimen strain

hardening which distributes the load over an ever-increasi ng area that

precludes continued temperature increases in the local area in which the

process initiated.

Magnesium Strengthening

Tensile yield strength data at low temperatures for three commercial aluminum-magnesium alloys (5083, 5154, and 5052) are presented

in Figure 6. The 5083-0 results represent the average data at 295, 76

and 4 K from the present study, whereas the 5154 and 5052 data were

previously unpublished [7]. In general, the temperature dependences are

very similar and little low-temperature strengthening is apparent.

Clearly, the addition of magnesium serves to increase the room-temperature

yield strength and this increase is maintained at low temperatures . In

Figure 7 the effects from magnesium additions to aluminum are plotted at

two test temperatures, 295 and 20 K. Again, the dependence of yield

strength on magnesium content at these two temperatures appears very

similar. This indicates that, contrary to nitrogen and carbon interstitial contributions to austenitic stainless steels, the addition of

magnesium serves to increase the long-range stress fields through solidsolution strengthening and does not affect the thermal contribution to

the flow strength. Data scatter in Figure 7 may represent the effects

of grain size or relate to lack of solute homogeneity.

191
SUMMARY

Tensile yield strength was measured for the 3.8-cm-thick plate of

aluminum alloy 5083-0 at temperatures ranging from 4 to 340 K. The

scatter of yield strength measurements was sizable, preventing an exact

determination of the temperature dependence. However, the following

results were obtained:

1. The stress-strain behavior of 5083-0 is complex: at 295 K

discontinuous yielding from dynamic strain aging occurs; at 4 K

discontinuous yielding from adiabatic specimen heating effects

occurs

.

2. Tensile yield strength is enhanced at room temperature by

dynamic strain aging from magnesium solute.

3. The strengthening contribution of magnesium essentially results

from its influence on long-range stress fields; little strengthening from short-range solute clustering or precipitation was

measurable from comparison of 295 and 20 K data.

ACKNOWLEDGMENTS

We thank D. Beard, Office of Fusion Energy, Department of Energy,

for his support and encouragement. D. Beekman and D. Baxter of the

National Bureau of Standards, Boulder, Colorado contributed by performing tensile tests, reducing data, and conducting metal lographic

analyses. Their assistance is gratefully acknowledged.

192
REFERENCES

1. Reed, R. P., Aluminum: A review of deformation properties of high

purity aluminum and dilute aluminum alloys. Cryogenics, 12 (1972)

259-291

.

2. Read, D. T. and Reed, R. P., Toughness, fatigue crack growth, and

tensile properties of three nitrogen-strengthened stainless steels

at cryogenic temperatures . The Metal Science of Stainless Steels ,

E. W. Collings and H. W. King, eds., American Institute of Mining,

Metallurgical, and Petroleum Engineers, New York (1979) pp. 92-121.

3. Metals Handbook , Vol . 2, Properties and Selection of Nonferrous

Alloys and Pure Metals, American Society for Metals, Metals Park,

Ohio (1 979).

4. Weston, W. F., Naimon, E. R., and Ledbetter, H. M., Low temperature

elastic properties of aluminum 5083-0 and four ferritic nickel

steels, in Properties of Materials for Liquefied Natural Gas Tankage ,

ASTM STP 579, American Society for Testing and Materials, Philadelphia

(1 975), pp. 297-420.

5. Morris, J. G., Dynamic strain aging in aluminum alloys, Mater. Sci

.

Eng. 13 (1974), 101-108.

6. Lloyd, D. J., The deformation of commercial aluminum-magnesium

alloys, Metal 1 . Trans. 1 1 A (1980), 1 287-1294.

7. Reed, R. P., Unpublished data. National Bureau of Standards, Boulder,

Colorado (1980).

8. Read, D. T. and Reed, R. P., Heating effects during tensile tests

of AISI 304L stainless steel at 4 K, Advances in Cryogenic Engineering ,

Vol. 26, Plenum Press, New York (1980), pp. 91-101.

193
LIST OF FIGURES

Figure 1

.

Figure 2.

Figure 3.

Figure 4.

Figure 5.

Figure 6.

Figure 7.

Three-dimensional metal lographic representation of aluminum

alloy 5083-0.

Tensile yield strength of 5083-0 aluminum alloy at low temperatures .

Variation of temperature dependence of tensile yield strength

of 5083-0 aluminum alloy from 4 K to the melting temperature.

Tensile true stress-true strain curves of 5083-0 aluminum

alloy at 295, 76, and 4 K.

Load and displacement controlled tensile stress-strain behavior at 4 K.

Tensile yield strength at low temperatures from aluminummagnesium alloys 5083, 5154, and 5052 in the annealed condition.

Influence of magnesium alloying additions in aluminum on

tensile yield strength at 295 and 20 K.

194
Figure 1

.

Three-dimensional metal 1 ographic representation of aluminum

alloy 5083-0.

195
YIELD

STRENGTH,

c

r

u

,

MPa

TEMPERATURE, K

Figure 2. Tensile yield strength of 5083-0 aluminum alloy at low temperatures .

196

YIELD

STRENGTH,

<rv

,

ksi
YIELD

STRENGTH,

cr.

(MPa)

0 100 200 300 400 500 600 700 800 900

TEMPERATURE (K )

Figure 3. Variation of temperature dependence of tensile yield strength

of 5083-0 aluminum alloy from 4 K to the melting temperature.

197

YIELD

STRENGTH/SHEAR

MODULUS,

cr,/ G

HO

3)
true

9

BOO

^true > ^

Figure 4. Tensile true stress-true strain curves of 5083-0 aluminum

* alloy at 295, 76, and 4 K.

198
Load Control

Displacement Control

Figure 5. Load and displacement controlled tensile stress-strai n behavior at 4 K.

199
YIELD

STRENGTH,

<rv

,

MPa

TEMPERATURE, K

Figure 6. Tensile yield strength at low temperatures from aluminummagnesium alloys 5083, 51 54, and 5052 in the annealed condition.

200

YIELD

STRENGTH,

<rv

,

ksi
TENSILE

YIELD

STRENGTH,

MPa

Figure 7. Influence of magnesium alloying additions in aluminum on tensile yield strength at 295 and 20 K.

201

EFFECTS OF CARBON AND NITROGEN ON THE ELASTIC CONSTANTS

OF STAINLESS STEEL 304

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Nine stainless-steel 304-type alloys were studied at room temperature. Carbon-plus-nitrogen contents of these alloys ranged from 0.067

to 0.325 weight percent (0.3 to 1.3 atomic percent). Five elastic

constants--longitudi nal modulus. Young's modulus, shear modulus, bulk

modulus, and Poisson's ratio--were determined by a pulse-echo ultrasonic

method. Within an experimental uncertainty of about 0.1 percent, all

the elastic constants remain unaffected by carbon and nitrogen additions .

203

Introduction

This study sought to determine how carbon and nitrogen alloying additions affect the elastic constants of stainless-steel 304-type alloys.

Both carbon and nitrogen increase strength of austenitic Fe-Cr-Ni

alloys. Nitrogen stabilizes the austenite, reducing the tendency at low

temperatures toward martensitic transformations. Carbon functions similarly, but higher c

r - rations cause welding sensitization. The increasing popularity of these alloys together with a dearth of reliable

physical -property measurements prompted the present study.

Being physical properties, elastic constants are relatively insensitive to details of internal structure related to thermomechanical

treatment. However, when determined with sufficient sensitivity, they

provide important parameters for indicating interatomic-force changes,

which may occur during alloying. Indeed, small changes in elastic

constants may indicate large changes in the nature of the interatomic

bondi ng

.

Experiment

Materials

Nine stainless-steel plates were purchased from the research laboratory of a U.S. steel manufacturer. All plates had nearly the same

base composition, which fell within the limits of ASTM specification

A240 for AI S

I -304 stainless steel, nominally Fe-18Cr-10Ni-l.5Mn-0.02P0.02S-0. 55Si-0.2Mo-0.2Cu weight percent. The carbon and nitrogen contents varied as shown in the mill chemical analyses in Table I. The

nine plates were produced from three 136-kg vacuum-induction-melted

heats, split with respect to carbon level, and teemed into 7.6 x 20 x

36 cm hot-topped cast-iron ingot molds. The ingots were reheated and

205
soaked at 1561 K (2350°F), hot-rolled into 2.54-cm plates, and air

cooled. The plates were finally annealed at 1332 ± 7 K (1937 ± 13°F)

for 1 h and water quenched. Hardness and grain size are also listed in

Table I, and mass densities are given in Table II.

Methods

Room-temperature sound velocities were determined by a pulse-echooverlap technique.^ Quartz piezoelectric crystals with fundamental

resonances between 4 and 7 MHz were cemented with phenyl salicylate to

one end of a specimen with flat and parallel surfaces. An x-cut transducer was used for longitudinal waves and an ac-cut for transverse

waves. The specimen in this case was a 1.59 x 1.59 x 1.27 cm parallelepiped with opposite faces ground flat and parallel within 2.5 ym.

Ultrasonic pulses 1 to 2 cycles long were launched into the specimen by

electrically exciting the transducer. The pulses propagated through the

specimen, reflected from the opposite face, and propagated back and

forth. The pulse echoes were detected by the transducer and displayed

on an oscilloscope equipped with a time delay and a microprocessor for

time-interval measurements. The sound velocity was computed by

v = 2£/t (1)

where £ denotes specimen length, and t the round-trip transit time. On

the oscilloscope, t was the time between adjacent echoes, the first and

second echoes usually being measured, and within these the time between

leading cycles. Elastic constants were computed from the general relations hi p

C = pv

2 (2)

206
where p denotes mass density. The usual engineering elastic constants

are related to the longitudinal and transverse sound velocities, v

£ and

v by

2 longitudinal modulus = = pv^ (3)

2 shear modulus = G = pv^ (4)

bulk modulus = B = C -(4/3)G (5)

Young's modulus = E = 3GB/(C^-G) (6)

Poisson's ratio = v = (E/2G) -1 (7)

= (l/2)(C r

2G)/(C

z -G)

Resul ts

Table II gives the principal results of the study: sound velocities

and elastic constants for nine stainless-steel type-304 alloys with different carbon and nitrogen contents. Table III shows for one case,

alloy 1, the directional variation of the ultrasonic wave velocities.

Table III implies also that ultrasonic velocities can be measured in

these alloys with an uncertainty of 0.05 percent or less.

Pi scussion

Results in Table II lead to the following conclusion: within our

measurement capability (0.1 percent or less) all elastic constants of

stainless-steel type-304 alloys remain unaffected by carbon-plus-nitrogen

additions up to 1.3 atomic percent. Various statistical analyses of the

results support this conclusion. And it can be seen quite simply by

comparing the results for specimens 1 and 9, which have the lowest and

207
highest carbon-plus-nitrogen additions: the two sets of elastic constants are nearly identical; neither volume-deformation nor sheardeformation elastic constants change significantly nor show any trend

versus composition. With the exception of specimen 4, all soundvelocity results fall within the one-standard-deviation range established for 304-type alloys by Ledbetter, Frederick, and Austin.^ And

specimen four is within 1.2 standard deviations.

Another strong reason for finding no alloying effect comes from

comparing the variation of B and G among the nine alloys. Interstitial

atoms, especially those in isotropic, octahedral positions, should

affect B more than G because they cannot move under a dilation force.

Yet, among the nine alloys B and G show the same variations from the

mean

.

Present results contrast sharply with observations reported for

2 ferritic iron. Speich, Schwoebie, and Leslie reported that carbon

added to alpha iron reduces elastic stiffnesses such as E and G by about

five percent per atomic percent carbon! Indeed, carbon is the most effective known alloying element in altering alpha iron's elastic constants. Quantitative results for nitrogen remain unreported, but one

would expect nitrogen to behave similarly to carbon because their atomic

diameters are similar. The essential distinction between the ferritic

and austenitic cases is the crystal structure. Body-centered-cubic

crystals are less accommodating to interstitials than are face-centeredcubic crystals. Carbon alloyed into bcc iron occupies octahedral interstices that have tetragonal symmetry, which produce a strong tetragonal

distortion; no similar distortion occurs in fee iron, where carbon

occupies larger octahedral interstices and the lattice distortion is

208
isotropic. This distortion reflects also in the much higher solubility

of carbon and nitrogen in fee iron compared with bcc.

Present results dispute those of a previous study. Laddach and

3 Schuler studied the influence of nitrogen on the elastic properties of

an Fe-21Cr-20Ni alloy containing 0.02 and 0.29 weight percent nitrogen.

They concluded that "nitrogen raises the modulus of elasticity as well

as the modulus of rigidity, regardless of the nitrogen being either in

interstitial solution or precipitated in the form of nitrides." No

reconciliation of this conclusion and that of the present study seems

apparent.

Conclusion

In the studied composition range, elastic constants are determined

by interactions among atoms occupying lattice sites; carbon and nitrogen

interstitial atoms play no significant part. Since elastic constants

are determined by the interatomic force constants, they also must be

independent of carbon and nitrogen content. Thus, the present result

has importance also for other force-constant-determined physical properties, such as thermal expansivity and specific heat.

Acknowledgments

The DoE supported part of this study. R. L. Tobler of NBS contributed the materials and the grain-size measurements. J. Dahnke of

NBS assisted in measurements.

References

1. H. M. Ledbetter, N. V. Frederick, and M. W. Austin, J. Appl

.

Phys. 51_, 305-309 (1980).

209
2. G. R. Speich, A. J. Schwoeble, and W. C. Leslie, Metall . Trans. _3,

2031-2037 (1972).

3. H. Laddach and P. Schuler, DEW-Techn. Ber. 1_3, 85-88 (1973).

210
Table

I.

Chemical

composition

(weight

percent),

hardness,

and

grain

size.

fl3 S- CD

LOCOOLOLOOOCOCXJO CONO^OCONOICOO^

CO

CO i CD r— C CD “O ^

ro U

m o

CtC

lo cr> lo cxj co o lo co co i

ro^-r^coco r^r^r^cocococoooco

3

CO

co o LO co cr» C\J LO CO f

cr> CXI O o •— »— •— CXI CXJ CXJ ro o o o o o o o O o

CT^ CTT cr> o o o o o o

ro ro ro CXJ CXI CXJ O O o •— — CXI (XI CXJ o o o o o o o o o

00 CO CT> r-* o LO

C\J LO CO ro LO CTT ro LO co o O o O o O o o o

o o o o o o o o o

o o ,

Oh o

CVI CXI o O o

LO LO o o CJh CXI CXI •— o o o

o LO

LO LO LO o o o

cr> LO

r— ir— o o o

o o o

CXI

CXI CXI CXJ o o o

o o o

r\ o CXJ LO LO

•— Oh cr» o CPi cr>

SCJ)

ro

CO oCO oCXJ

.— cxjrO'^-Lntotxoocn

211
Table II. Mass density, sound velocities, and elastic constants of nine

carbon- pi us-ni trogen-al 1 oyed stainless-steel 304 alloys.

Densi ty

(g/cm 3

)

v ,, (cm/ys)

v

t (cm/us)

C

£ (GPa)

E (GPa)

G

(GPa)

B (GPa)

V

1 7.944 0.5751 0.3140 262.7 201 .7 78.3 1 58 .

3

0.288

2 7.937 0.5757 0.3139 263.1 201 .5 78.2 158.8 0.288

3 7.936 0.5751 0.3142 262.5 201 .7 78.3 158.0 0.287

4 7.935 0.5743 0.3155 261 .7 202.8 79.0 156.4 0.284

5 7.927 0.5747 0.31 50 261 .8 202.2 78.7 156.9 0.285

6 7.923 0.5748 0.3144 261 .8 201 .5 78.3 157.3 0.287

7 7.869 0.5769 0.3141 261 .9 200.2 77.6 158.4 0.289

8 7.883 0.5764 0.3135 261 .9 199.9 77.5 158.6 0.290

9 7.900 0.5766 0.3146 262.6 201 .4 78.2 158.4 0.288

Mean 0.5755 0.3144 262.2 201 .4 78.2 157.9 0.287

Std

.

dev

.

0.0009 0.0006 0.5 0.9 0.5 0.8 0.002

Pet

.

std. error 0.2 0.2 0.2 0.4 0.6 0.5 0.7

212
Table III. Directional variation of ultrasonic wave velocities in

specimen 1, where z direction is perpendicular to rolling

pi ane

.

Propagation

di rection

Polarization

di rection (cm/us)

v

t (cm/ys)

X X 0.5746

X y 0.3140

X z 0.3143

y y 0.5753

y X 0.3139

y z 0.3138

z z 0.5753

z X 0.3143

z y 0.3140

Mean 0.5751 0.3141

Std. dev. 0.0004 0.0002

Pet. std. error 0.07 0.07

213

LOW-TEMPERATURE VARIABILITY OF STAINLESS-STEEL-304 ELASTIC CONSTANTS

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

By measuring ultrasonic velocities, the elastic-constant behavior of four

stainless-steel 304-type alloys was determined between ambient and liquidhelium temperatures. Despite their equivalence near room temperature, results

show that cooling through the magnetic transition produces slightly different

elastic constants in different alloys at low temperatures.

215

INTRODUCTION

Recently the author 1 collaborated in studying the room- temperature

variability of stainless-steel 304-type elastic constants: Young's modulus,

shear modulus, bulk modulus, Poisson's ratio. For twenty samples acquired

randomly, the lot-to-lot variation was surprisingly small: 0.9 percent for

Young's modulus, for example. Thus, for this commercial alloy, chemical -composition differences and thermomechanical -processing differences affect the elastic constants practically negligibly.

Because of many cryogenic applications for 300-series austenitic stainless

steels, the author2 recently determined accurately the low- temperature elastic constants of three such alloys: 304, 310, 316. At low temperatures all

three alloys undergo a magnetic transition, which affects the elastic constants. The 304-alloy results given in that study were representati ve rather

than comprehensive. Different 304 alloys actually behave slightly differently

at low temperatures.

The present study determined approximately the low- temperature variability

of stainless-steel -304 elastic constants. The longitudinal sound velocity was

measured between 295 and 4 K for four steels: one regular 304 alloy, two lowcarbon alloys, and one nitrogen-enriched alloy. The author3 showed recently

that interstitial alloying elements such as carbon and nitrogen do not affect

the room- temperature elastic constants of 304-type stainless steels. The

longitudinal modulus manifests both volume and shear effects and is therefore

sensitive to any change in the elastic constants.

217
EXPERIMENTS

Low- temperature ultrasonic velocities were determined similarly to roomtemperature measurements made in our laboratory by a pulse-echo-superposition

method.^ >3 The notable difference was that for low temperatures the transducer-specimen bond consisted of stopcock grease rather than phenyl salicylate.

The specimen was inserted into a probe, 4 which was inserted into the ullage

of a helium dewar. Measurements were made at 5-K intervals during continuous

cooling or warming at rates of 1 to 2 K per minute. Temperature was monitored

with a chromel -constantan thermocouple placed close to the specimen.

Table I shows the chemical compositions, hardnesses, and mass densities of

the studied alloys.

RESULTS

Figure 1 shows the experimental results graphically: relative longitudinal modulus versus temperature for six alloys. To determine the slope dC/dT

more accurately, phenyl salicylate bonds were used in two cases and cooled only

to 245 K. Table II gives quantitative character!' sties of the cooling curves

(which reproduced during subsequent reheating, showing no hysteresis). "Max"

indicates the maximum value of the C ( T) /C ( 295 ) curve, "min" indicates the

minimum value at low temperatures, and A the difference between the two. The

transition temperature, Tt

, corresponds to A/2. The slope dC/dT was determined by a 1 inear-regression fit to eleven points between 245 and 295 K, and

r^ is the coefficient of determination.

218
DISCUSSION

Elastic constants reflect the magnetic-interaction energy of a system

according to

C

ijkl

=

3 2

U/3Tl

1-j3nkl (1)

where C-fj^i denotes the fourth-rank elastic-stiffness tensor, U denotes

internal -energy density (including magnetic energy), and n-jj denotes Brugger

strains. Four types of magnetic energy were discussed by Ledbetter and

Col lings5 in a study of Fe-Cr-Ni-Mn steels.

A priori , one expects strong magnetic interactions in these steels because

they contain four magnetic elements: iron, chromium, nickel, and manganese.

Fundamental calculations by Moruzzi , Janak, and Williams5 show that for these

elements, especially Mn, magnetic energy affects properties: cohesive energy

is reduced, bulk modulus reduced, and volume increased. At lower temperatures,

because lattice vibrations are reduced and because magnetic atoms move closer

together because of thermal contraction, magnetic interactions become stronger.

Stronger magnetic interactions manifest themselves in physical -property

changes.

The existence and complexity of low- temperature magnetic transitions in

304-type alloys have been discussed especially by Col lings and Hart. 7 They

view austenitic stainless steels as completely paramagnetic at high temperatures. At some lower temperature, clustered ferromagnetic ordering leads to

219
superparamagnetism; at still lower temperatures , the paramagnetic matrix undergoes spin-glass condensation. Interactions among spin clusters and with the

matrix lead to further possible complexities. In principle, all these magnetic

transitions affect the elastic constants.

The reversible two-percent elastic-constant softening during cooling shown

in Fig. 1 near T = 47 K results from a magnetic transition, probably the spinglass condensation.

Why should spin-glass condensation cause elastic softening? Landau and

Lifshitz8 proved for second-order phase transitions that "compressibility

drops for a transition from an unsymmetrical to a symmetrical phase." Thus,

the bulk modulus drops for a transition from a symmetrical (paramagnetic) phase

to an unsymmetrical (magnetic) phase. Magnetic in this case means spin-glass.

Spin-glass properties arise, apparently, from magnetic-moment interactions. Principal among these is the s-d or Rudermann-Kittel-Kasuya-Yosida

(RKKY) indirect-exchange interaction. The energy function for this spin system

is

( 2 )

where (3)

J ( R) = A cos 2k

f

R/(2k

f

R)

3

(4)

220
where the spin positions are random variables, kf is the fermi-level

wavevector of the host conduction electrons, and J ( 0) = 0. Interaction (4)

alternates in sign and is long range, representing al ternatingly ferromagnetic

and anti ferromagnetic spin coupling. This type of magnetic interaction can

become important when the usual S -j • S

j

direct exchange becomes insignificant,

that is, when the magnetic ions are separated so widely that their overlap

integral is zero. The basic idea in RKKY exchange is that an s conduction

electron interacts with a localized d electron through a direct-exchange interaction S

s -Sf, travels through the lattice, and then has a second such interaction S

s, Sj. This amounts to an effective (indirect) interaction

between the two localized d electrons at "Rj and Rj.

The anomalous elastic-constant behavior shown in Fig. 1 relates in the

most general sense to magneto-elastic energy, to a mechanism that couples a

crystal's elastic vibrations to its magnetic modes. Here we are especially

concerned with the magneto-elastic energy arising from the strain dependence of

the JjjSj'Sj indirect exchange interaction. From Hausch and Warlimont,^ who

used a molecular-field approximation, and for simplicity using Voigt averaging,

it follows that the magnetic-exchange-energy contributions to the bulk modulus,

B, and the shear modulus, G, for the face-centered-cubic first neighbor-only

case are:

B = -4N( 1 - J" - 2r J')

ex T.J. o o

(5)

221
and Gex^^(1 -T ( 6 ) t

> (r

o

J

" + 4r

o

J '»

where N = atoms/volume, Tt

= transition temperature, r0= equilibrium

atomic spacing, and J' and J" are the first and second derivatives of J with

respect to r evaluated at r = r 0. Since J(r) is not known a priori for 3d

metals and alloys, J' and J" must be determined experimentally. In principle,

J' and J" can be estimated from measurements made in the present study. However, the many approximations and the alloy's complexity render this approach

dubious. Equations (5) and (6) are set out mainly to give the flavor of a

final explanation, which awaits an ab initio relationship for J(r), discussions

of which Dehlinger1 ^ gave for alloys.

ACKNOWLEDGMENT

The Department of Energy supported this study. Several people contributed

experimentally: M. W. Austin, J. Dahnke, and G. Maerz.

REFERENCES

1. H. M. Ledbetter, N. F. Frederick, and M. W. Austin, J. Appl . Rhys. 5_1,

305-309 (1980).

2. H. M. Ledbetter, J. Appl. Phys., forthcoming.

222
3. H. M. Ledbetter, unpublished research.

4. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics 246-249

(1974).

5. H. M. Ledbetter and E. W. Collings, in The Metal Science of Stainless

Steel

s

(Met. Soc. AIME, New York, 1979), pp. 22-40.

6. V. L. Moruzzi, J. F. Janak, and A. R. Williams, Calculated Electronic

Properties of Metals (Pergamon, New York, 1978), p. 6.

7. E. W. Collings and S. C. Hart, Cryogenics T9, 521-530 (1960).

8. L. D. Landau and E. M. Lifshitz, Statistical Physics (Pergamon, London,

1959), p. 438.

9. G. Hausch and H. Warlimont, Z. Metal lkd. 63, 547-552 (1972).

10.

U. Dehlinger, Z. Metallkd. 28, 116 (1936); 28, 194 (1936); 29, 388 (1937).

Reviewed in F. Seitz, The Modern Theory of Solids (McGraw Hill, New York,

1940), pp. 623-626.

223
Table

I.

Chemical

composition

(weight

percent)

and

other

properties

So—

-MOO

CO -r- £ oo CO ['- 0 0 1—

co co

u CM CO CM ^3- CO 1—1

HJ c\ Ch co CTi <T> 00 00 2 (1) D) • • • ® 9 • o r-. r^- r— f- r~-

CO

CO

co i

O) r— C <D r- C\J CM co LO r-- -O 5 oo co cn i"- r- r~-

-ic to o

n: o

cc —

co o 0 C_> 9 0

CO0 > 9 O 1

CM o CM LO s • 9 O O

co C- o C\J O C_3 9 9 0 O -—i

CO CM Z 0 rH • 9 0 0 0 r- CO <vJ- pH rH rH 1 I CM

CO o 0 0 0 O • 9 • • 9 o 0 0 0 O

C\J 0 LO CTi CM CM CM CM rH

o_ o OOO 0 • • • • 9 o O O O 0

CO rH

O co

CM CM CO CO o O 0 0 0 o • a • 9 9 o 0 0 O 0

CM CT> rH 00

CO CO LO CO CO

00 • • • 9 9 o 0 O O O

LO CM CO CO c CO CO co s • 9 • 9 9 rH rH O rH rH

co LO LO <Ti r— r-1

•r- • • 9 9 « 0 co 00 cn rH*3” 0 LO

i- CO CO LO co o • a • • 9 CO co co co co

1-1 rH rH T—(

pH CM CO LO co

224
Table II. Results of low- temperature measurements of relative longitudinal modulus

max min A Tt

(K)

-dC/dT

(10 7 N/m2 K)

r

2

1 1.0442 1.0248 0.0194 46 6.62 0.9909

2 6.55 0.9944

3 1.0436 1.0207 0.0229 45 6.97 0.9910

4 1.0421 1.0147 0.0274 51 7.37 0.9917

5 6.04 0.9978

6 1.0429 1.0184 0.0245 47 6.38 0.9912

225
(T)/C,

(295)

1.05 T

o

TEMPERATURE, K

Fig. 1 Relative longitudinal modulus versus temperature for six stainlesssteel type-304 alloys. This modulus reflects both dilation and shear deformation modes. As described in text, anomalous decrease at near 47 K arises from a paramagnetic-magnetic transition, probably spin-glass, which

is usually described by an RKKY indirect exchange energy.
PREDICTED SINGLE-CRYSTAL ELASTIC CONSTANTS OF STAINLESS STEEL 304

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

The three single-crystal elastic constants--^, C-^, C^--of

stainless steel 304 are estimated from the polycrystalline bulk and

shear moduli together with an empirical C^/C^ value, which is discussed theoretically. The estimate involves a reverse Kroner method for

relating single crystal and polycrystal 1 ine elastic constants.

227

A solid's engineering elastic constants, such as Young's modulus

and Poisson's ratio, serve many practical purposes. Among these are

calculations of load-displacement, thermoelastic stress, elastic stability (buckling), residual stress, and fracture-toughness pi ane-stressto-plane-strain conversions.

A solid's single crystal (anisotropic) elastic constants, the C..'s

J and the S..'s, provide additional and more fundamental information than

J do the engineering (quasi -isotropic ) elastic constants. The C..'s are

' J essential, for example, for theory of point defects, anisotropic dislocation interactions, elastic-energy calculations , force-constant-related

properties such as lattice vibrations, accurate determination of Debye

characteristic temperatures, correlations with electronic theory of

solids, and correlations with other physical properties.

Stainless steels form a key component of materials technology, and

their basic properties are now being studied by methods applied formerly

to elements or to simple binary alloys. Elastic constants comprise one

subset of basic mechanical -physical properties, which must be known and

understood before the metal physics of stainless steels can be carried

very far.

Three existing experimental studies [1-3] of the single-crystal

elastic constants of Fe-Cr-Ni alloys omit the Fe-18Cr-8Ni alloy, which

is of most interest because of large tonnages and multifarious applications and because it provides a base material for developing many other

special purpose stainless steels. In principle, these three studies on

four alloys--Fe-12Cr-18Ni , Fe-18Cr-l 2Ni , Fe-18Cr-l 4Ni , Fe-18Cr-l 9Ni

should permit by linear-regression analysis the determination of the

singl e-crystal elastic constants of Fe-18Cr-8Ni. But, this analysis

229
fails. Reasons for this probably include: a small sample set (four

alloys for three independent coefficients), poor distribution of sample

chemical compositions, experimental errors, nonlinear dependence on

composition, and possible nonadditive, interactive composition effects.

The present study predicts the si ngl e-crystal elastic constants of

Fe-18Cr-8Ni. The prediction depends on three ingredients: one observational, one theoretical, and one empirical, which has some theoretical

basis. Observational ly, Ledbetter, Frederick, and Austin [4] measured

the polycrystal 1 i ne elastic constants of twenty randomly acquired Fe18Cr-8Ni alloys with the results that the bulk modulus, B, and the shear

modulus, G, are well defined: B = 1.582 and G = 0.774* 1 0^ N/m^ . Theoretically, Kroner [5] showed how the polycrystal 1 ine elastic constants

(two independent, e.g., B and G) can be calculated from the singlecrystal elastic constants (three independent for the cubic-symmetry

case: C^, C-^, and C^). Empirically, as shown in table 1 for Fe, Ni,

and the four Fe-Cr-Ni alloys described above, the dimensionless ratio y

^12^11 1S surprisingly constant for these materials: 0.635 ± 0.031.

In theory, y varies between zero and one. Both the upper and lower

limits correspond to a mechanical instability. For the cubic elements,

in practice y varies between 0.12 (diamond) and 0.85 (gold).

Before considering the above case, we consider a simple, familiar

model: central forces between atoms, which require that = C^.

Table 1 shows this assumption is not too bad for these materials. We

will treat the results as approximate guidelines to a better calculation. Since we are computing three elastic constants (C-|-|, C.^, C^)

from two (B and G), we need another relationship among the C^.'s. For

central interatomic forces, this relationship is

230
C

1 2

C 44 ( 1 )

For all cubic crystals, the bulk modulus relates to the C..'s according

J to

B "

3

^ C

ll + 2C 12^ ( 2 )

Following Kroner, the shear modulus relates to the C..'s by

3 *3

G

2 + aG2 + BG + y

= 0 (3)

where

a = (5C n

+ 4C

12 )/8 (4)

3 = -C

44

(7C n

- 4C

12 )/8 (5)

and

Y

= -C

44

(C n

- C

12

)(C n

+ 2C

12 )/8 (6)

Solving eqs . (1) through (3) simultaneously gives the central -force (cf)

results for Fe-18Cr-8Ni shown in table 1. All three C - .

1 s and the bulk

J modulus fall within the ranges established by the four alloys. Both

C' = (C-|i-C-| 2 )/2 and A = C 44

/ C

' are outside these bounds. Thus, when

fitted to observed values of B and G, a central -force model for stainless steel is not bad and fails mainly for the C' elastic constant.

We turn now to a more realistic noncentral -force model where C^

2 /C^

0.635. (The basis for this is described below.) Simultaneously solving

eqs. (2) and (3) subject to this condition yields the results shown in

table 1 in the bottom line. When combined with previous observations,

these new theoretical results show how Ni affects elastic constants in

the Fe-18Cr-8Ni region. Increasing Ni decreases , C-^, B, and C

1 ,

231
but leaves relatively unaffected, perhaps slightly increased; Zener

anisotropy increases slightly, indicating a possible tendency toward

instability. Except for C' and A, results of Salmutter and Stangler [1]

suggest that Cr produces opposite effects from Ni

.

Finally, it remains to try to understand why y

= 0.635 for these

alloys. If the two-body, central -force interatomic potential is denoted

4> ( r ) , then the Brugger elastic constants, which are fourth-rank tensors,

are given at T = 0 K by:

c =

i j kl 8n • • 3n 1

1

J ij kl

(7)

=

2W ^ D2^ r ) r

i r

j r

k r l^ r = r

0 ’

where D<f>(r) denotes (1/r )[3cf>(r)/3r] . The energy density U of the

crystal is obtained by a sum over two-body atom-atom interaction energies :

U = (l/2V°)z<f>(r), (8)

where V° is the undeformed atomic volume. The n

n

- • are components of the

Lagrangean elastic strain matrix

n=}(J t J-I), (9)

where I is the 3x3 identity matrix, t denotes transposition, and the

deformation matrix J is defined by

r = Jr° , (10)

232
where r° and r are the interatomic spacings in the initial state and in

the deformed state, respectively . Details concerning this approach to

elastic constants can be found elsewhere [9].

Face-centered cubic crystals have twelve nearest neighbors at

<^±1 , ±1, o)>a/2, and six second-nearest neighbors at <±1 , 0, Cj>a, where a

is the unit-cell dimension. In this case one obtains from eq. (7):

C

1

1

= 2 [8(|-)^D 2

<j> (—) + 2a^D

2

<j)(a) + ...] (11)

/2

and

C-|2 = ^44

= 2^ [ 4 (|0

^ D^<j) (—) + ...], (12)

where ... indicates contributions from pairs farther than second neighbors. Clearly, summing over only nearest neighbors gives y

= 0.500 for

the f.c.c. case. The effects of second-nearest neighbors depend on

2 whether D ( r ) is already negative at r = a. If so, then y is increased

because is decreased while remains unaffected. Equation (7)

represents both structure-dependent and volume-dependent contributions

to the C..'s. But there may be other, purely volume-dependent terms in

* J the energy, often represented as

U( V) = AV

n (13)

where V denotes volume and A and n are positive adjustable parameters.

It follows simply that these terms (denoted by asterisks) contribute to

the C. . 's as follows

:

J

*

C

11

= nA(n-2) = n^A-2nA (14)

*

rL

1

2

= n

2 A (15)

*

C (16) 44

= nA

233
Thomas [10] showed that for f.c.c. lattices

-nA = /2a(3<j>/3r) (17)

evaluated at the nearest-neighbor spacing. Since 3<f>/3r is negative

there, nA is positive. (If the volume-dependent energy arises, for

example, entirely from the electron gas, then n = -2/3 in a Sommerfeld

model.) Thus, C-^* exceeds C-j-|* and y is increased above 0.50 by

volume-dependent energy terms. For cubic metals, only four (Cr, Ir, Mo,

W) have C-^/C-ji ratios less than 0.5, and only one of these is f.c.c.

Values exceeding 0.50 are expected for metals because their electron gas

contributes a volume-dependent energy. The near constancy of y among

Fe, Ni, and these Fe-Cr-Ni alloys suggests that for some purposes the

potential -energy function of d-band transition metals can be treated

similarly to the simpler sp-band metals: a sum of pair potentials and a

purely volume-dependent energy term.

ACKNOWLEDGMENT

This study was supported partly by the DoE Office of Fusion Energy

and partly by the NBS Office of Standard Reference Data.

REFERENCES

[1] K. Salmutter and F. Stangler, Z. Metal Ikd. 51_, 1-4 (1960).

[2] M. Kikuchi, Trans. Jap. Inst. Met. 1_2, 417-421 (1 971 ).

[3] M. C. Mangalick and N. F. Fiore, Trans. Metal! . Soc. AIME 242 , 363

2364 (1968).

[4] H. M. Ledbetter, N. V. Frederick, and M. F. Austin, J. Appl . Phys . 51_,

305-309 (1980).

234
[5]

E. Kroner, Z. Phys. 1_51_, 504-518 (1958).

[6] R. A. Johnson, Phys. Rev. 145 , 423-433 (1966).

[7] D. Diesburg, Ph.D. Thesis, Iowa State University (1971).

[8] G. A. Alers, J. R. Neighbors, and H. Sato, J. Phys. Chem. Solids,

13., 40-55 (1960).

[9] R. A. Johnson, Phys. Rev. B 6^, 2094-21 00 (1972).

[10] J. F. Thomas, Scripta Metal 1. 5., 787-790 (1 971 ).

235
Table 1

Single-Crystal Elastic Constants of Several

Face-Centered-Cubic Materials

Material C

11 C1

2

C 44

lO^N/m 2

B r A y Ref.

Fe 2.51 1 .57 1 .12 1 .88 0.47 2.38 0.625 6

Fe 2.760 1 .735 1 .363 2.077 0.513 2.41 0.629 7

Ni 2.516 1 .544 1 .220 1 .868 0.486 2.51 0.614 8

Fe-1 2Cr-l 8Ni 2.332 1 .626 1 .225 1 .862 0.353 3.48 0.697 1

Fe-1 8Cr-l 2Ni 1 .912 1 .179 1 .386 1 .423 0.367 3.78 0.617 2

Fe-18Cr-14Ni 1.98 1 .25 1 .22 1 .49 0.365 3.34 0.631 3

Fe-18Cr-l 9Ni 1 .91 1 .19 1 .24 1 .43 0.360 3.44 0.623 3

Fe-18Cr-8Ni (cf) 1.99 1 .38 1 .38 1 .58 0.306 4.50 0.692 Present

Fe-1 8Cr-8Ni 2.09 1 .33 1 .21 1 .58 0.382 3.17 0.635 Present

236
ELASTIC-CONSTANT/TEMPERATURE BEHAVIOR OF THREE HARDENED MARAGING STEELS

H. M. Ledbetter and M. W. Austin

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Elastic constants of three maraging steels were determined by

measuring ultrasonic velocities. Annealed steels show slightly lower

bulk moduli and considerably lower shear moduli than hardened steels.

All the elastic constants -- Young's modulus, shear modulus, bulk

modulus, and Poisson's ratio -- show regular temperature behavior

between 76 K and 400 K. Young's and shear moduli increase with increasing yield strength, but the bulk modulus and Poisson's ratio are

relatively unchanged. Elastic anisotropy is quite small.

237

INTRODUCTION

Elastic constants find many uses. In technological materials they

assume importance whenever problems arise concerning deflections, elastic

stability (buckling), excessive elastic deformation (jamming), thermoelastic

stress, residual stress, fracture, and so on. Ratios of elastic modulus

to density or to thermal conductivity provide well-known structuraldesign parameters. Elastic constants correlate empirically with many

practical properties such as strength, hardness, and wear. Because they

relate intimately to both the interatomic potential and the latticevibration spectrum, elastic constants relate in turn to many fundamental

solid-state phenomena such as diffusion, theoretical strength, and

lattice-defect energies. Elastic constants appear as coefficients in

thermodynamic equations of state where their relationships to free

energy, specific heat, and thermal expansivity become especially clear.

Since they can be measured accurately (typically within 1 percent or

5 less) and small changes can be detected easily (typically 1 in 10 or

better), elastic constants are useful physical properties for characterizing

and comparing materials and for detecting changes in material properties.

In the present case, the variation of elastic constants with composition

and microstructure is of special interest.

Ledbetter and Read^ recently reported the elastic-constant behavior

of an annealed 300-grade maraging steel between 4 and 300 K. [The 300-

grade steel has a yield stress near 2.07 GPa (300,000 psi).] They

reported four elastic constants -- Young's modulus, shear modulus, bulk

modulus, and Poisson's ratio. Besides corresponding closely to the

elastic constants of an Fe- 1 8N i alloy, all these constants showed regular

temperature behavior.

239
The present report describes a similar, extended study. Three

maraging steels -- 200-grade, 250-grade, and 300-grade, both annealed

and hardened -- were studied between 76 and 400 K. Elastic anisotropy

was studied at room temperature. Despite their relatively poor lowtemperature fracture toughness in the hardened condition, these steels

occur in many cryogenic applications because of their exceptionally high

strength. Because they are relatively new structural alloys, their

elastic properties have received little attention.

EXPERIMENT

Tables I and II characterize the alloys obtained from a commercial

source as 2.7-cm and 4.2-cm diameter bars. The pycnometric mass-density

determination used distilled water as a standard.

Sound velocities were determined on two sizes of cubes -- 1.3 cm

and 1.6 cm -- whose opposite faces were flat and parallel within 3 pm.

Pulse-echo and pulse-echo-superposition techniques for room-temperature

2 3

and low-temperature measurements, respectively, are described elsewhere. ’

Measurements between 300 and 400 K were done in a stirred, heated mineraloil bath using the oil as the transducer-specimen couplant; a mercuryin-glass thermometer indicated the temperature.

Elastic constants were obtained from longitudinal and transverse

sound velocities, v^ and v^, by the relationships:

240
2 longitudinal modulus = C

£ = pv£ (1)

2 shear modulus = G = pv

t (2)

4

bulk modulus = B = C

£

-

^-G (3)

Young's modulus = 9GB/(G + 3B) (4)

Poisson's ratio = (E/2G) - 1 (5)

where p denotes mass density.

RESULTS

Table III gives, for the hardened steels, the room-temperature

elastic constants and their directional variation. Table IV gives the

averaged (over three orthogonal directions) room-temperature elastic

constants for the annealed specimens. Table V and Figs. 1-5 give the

temperature dependence of the elastic constants. Curves in these figures

4

represent best fits of the data to Varshni's relationship:

C = C° - s/(e t/T

- 1) (6)

which contains three adjustable parameters: C° is the zero-temperature

elastic constant, -s/t is the high-temperature slope dC/dT, and t is an

approximate Einstein temperature. Figures 1-5 contain only the 76 to

300 K results; the 300 to 400 K data form linear extensions of the 76 to

300 K curves.

241
DISCUSSION

In their annealed ( Fe-Ni-martensite) condition, these alloys show

decreasing E and G and increasing B and v upon progressing from lower to

higher yield strength (see Table IV). Since only Mo content increases

systemati cally in this series and its bulk modulus exceeds considerably

5 that of Fe, Ni, and Co, it may cause these property variations. Takeuchi's

elastic-constant results confirm that Mo additions decrease Fe's shear

modulus. Molybdenum effects on the bulk modulus remain uncertain.

Ledbetter and Read"' found that, except for the bulk modulus, an annealed

300-grade maraging steel has elastic constants similar to those of

F e - 1 8 N i . The present study confirms this.

Hardening increases Young's modulus, shear modulus, longitudinal

modulus, and bulk modulus but decreases Poisson's ratio (see Tables III

and IV). These changes increase upon progressing from 200-grade to

250-grade to 300-grade. These elastic-constant changes suggest their

possible use as an indicator of the precipitation process.

The elastic-constant temperature behavior is hardly remarkable. It

shows all the features of regular behavior: tendency toward zero slope

at zero temperature, as required by the third law of thermodynamics

;

continuous decrease with increasing temperature, consistent with the

softening of interatomic bonding forces caused by increased thermal

vibrations; and a linear slope at high temperatures . High-temperature

linearity is consistent with the theoretical result of Leibfried and

Ludwig that the elastic constant depends on temperature according to

C(T) = C(0) (1 -DU) (7)

where U denotes mean energy and D is a parameter that depends on the

type of crystal and the type of oscillator model used (Einstein, Debye,

etc.). At high temperatures U varies as kT.

242
ACKNOWLEDGMENT

This study was supported by NASA Langley Research Center and by the

DoE Office of Fusion Energy.

243
REFERENCES

1. H. M. Ledbetter and D. T. Read, Metal! . Trans. A 8A, 1805-1808

(1 977).

2. E. R. Naimon, W. F. Weston, and H. M. Ledbetter, Cryogenics ]_4,

246-250 (1 974).

3. H. M. Ledbetter, N. F. Frederick, and M. W. Austin, J. Appl . Phys.

51_, 305-309 (1 980).

4. Y. P. Varshni, Phys. Rev. 2 , 3952-3955, (1970).

5. S. Takeuchi, J. Phys. Soc. Jpn. 27_, 929-940 (1 969).

6.. G. Leibfried and W. Ludwig, in Solid State Physics, Volume 12

(Academic, New York, 1961), pp. 275-444.

244
Table

I.

Chemical

compositions

of

studied

steels,

weight

percentage,

balance

Fe.

245
Table II. Metallurgical characterization of studied alloys.

Densi ty

(g/cm3

)

Hardness

(Rockwell C)

Hardening Heat Treatment

Annealed Hardened Annealed Hardened

200-grade 8.068 8.030 28 43 489°C

250-grade 8.076 8.066 29 51 same

300-grade 8.087 8.087 31 54 same

246
Table III. Room-temperature elastic constants

of three hardened maraging steels.

B E G V

(10

11

N/m

2

) (10

11 N/m 2

) (10

11 N/m

2

)

200-grade x

a

1 .654 1 .853 0.706 0.313

y 1 .657 1 .848 0.703 0.314

z 1 .656 1 .852 0.705 0.314

250-grade X 1 .666 1 .885 0.719 0.311

y 1 .664 1 .884 0.718 0.311

z 1 .663 1 .888 0.720 0.310

300-grade X 1 .680 1 .902 0.725 0.311

y 1 .681 1 .899 0.724 0.312

z 1 .686 1 .899 0.724 0.312

a

z = bar axis; x and y are perpendicular to z.

247
Table IV. Room-temperature elastic constants of three annealed maraging

steel s

.

B

(10

11 N/m2

)

E

(10

11 N/m

2

)

G

(10

11 N/m2

)

V

200-grade 1.655 1.773 0.671 0.321

250-grade 1.656 1.749 0.661 0.324

300-grade 1.664 1.729 0.652 0.327

248
Table V. Elastic constants of three hardened

maraging steels at selected temperatures

.

T(K)

(1

0

1

1

N/m

2

) (10

11 N/m

2

) (10

11 N/m

2

)

200-grade

400 1.63 1.794 0.681 0.317

350 1.643 1 .822 0.693 0.315

300 1.655 1.849 0.704 0.314

250 1.667 1.875 0.714 0.313

200 1.678 1.900 0.724 0.311

150 1 .687 1.920 0.733 0.310

100 1 .694 1.935 0.739 0.310

75 1.696 1.938 0.740 0.309

250-grade

400 1.645 1 .826 0.695 0.315

350 1.654 1 .855 0.706 0.313

300 1.664 1.883 0.718 0.311

250 1.673 1 .910 0.729 0.310

200 1.682 1.935 0.740 0.308

150 1 .691 1.956 0.748 0.307

100 1.698 1 .971 0.754 0.307

75 1 .701 1.976 0.756 0.306

300-grade

400 1.664 1 .841 0.700 0.316

350 1 .674 1.870 0.712 0.314

300 1.684 1.898 0.723 0.312

250 1.694 1.925 0.734 0.311

200 1.704 1.949 0.744 0.309

150 1.713 1 .971 0.753 0.308

75 1.72 1 .986 0.759 0.308

75 1 .722 1.990 0.761 0.307

249
LIST OF FIGURES

Fig. 1

Fig. 2

Fig. 3

Fig. 4

Fig. 5

Temperature variation of longitudinal modulus of three

hardened maraging steels.

Temperature variation of shear modulus of three hardened maraging

steel s

.

Temperature variation of Young's modulus of three hardened maraging

steel s

.

Temperature variation of bulk modulus of three hardened maraging

steels.

Temperature variation of Poisson's ratio of three hardened maraging

steel s

.

250
o

G

m

Q

<3

nj

Q

Q

-a

ajc

a> x>

s03

CL) CD

£_ .c

4->

C/1 13

-a

03 p— • i- C/1 03 c

> a)

a>

ai +-> S_ 00

13

+-> CD

03 C

S_ -r- aj o>

Q. 03 E S- a) 03 i- E

C\J G3 ao LD OJ a 03 • r- 0- ID ID CD CD CD in o>

• • • > • • • • •1— OJ OJ OJ OJ OJ OJ OJ OJ OJ Li_

(2^^/n IW0D smnaow lUNiamiDNcn

251

TEMPERATURE

CK)
<3

U

d

D

1-d

d

LJ

CLZ

Ld

-o

cuc

cu •a

i.

03

CU

a>s_

4-O

01

13

-ao

s-03

CU

4-O

+-> 03

• i— • S- 00

03 i

> CU

d)

CU +J

5- 00 3

+-> 03

03 C

-r- CU CT3 Q. 03 E SCU 03

C\J

CT>

CD

rv rv. CDrv mrv

•<r [V

(D

rv

OJ

rv Q

rv

(2^uj/n iW0i) smnaow dd3HS

252
<sm T3

cu

CU -a

src

CL) O)

sJZZ

4-> o

COD

IS 1—

IS 13 PJ ~oo

E w CO u CD cr c

D D O

1- >- CE 4- CK O

U C

CL o

21 •f

+-> Ld fa Q h- •1— • IS S- CO

CO r— > cu

CU

CO

cu

s-=j

+-> a>

CO C

5- *r- CU CD Q. ra E s- <u ra

Q

CD

OJ O GO CD 'sr CM (S GO CD Q IS) cn cn 01 CD CD 00 GO GO

• • • • • • • • • • OJ OJ —

*

rH

( 2 vUJ /N IN0D smnaow s.DNnox

253
CD

CD

ra

s~

<T3

*o

CDc

cu a

s_

03

CU

cu

s_

o

00

=3 -ao

rs .o

c

o

4-

>

<T3 •r—s_

ra>

<u

5-

Z5

<T3 S- 00

CD r—

CL CU E CD CU +-> I— CO

rrPw ru

IN. S3

PCDCD

CO

CD CO

(2v-uj/n IW0D smnaow >nne

254
oi Ida s Nossiod

-a

CDc

CD T3

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255

TEMPERATURE BEHAVIOR OF YOUNG'S MODULI OF FORTY ENGINEERING ALLOYS

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Young 1 s-modul us and temperature data are collected graphically and

tabularly for forty alloys that have technological applications. Alloy

base metals include: aluminum, copper, iron, and nickel. Sources of

data are: handbooks, original research at NBS , and the scientificengineering literature. The temperature range 0 to 590 K (-460 to

600° F) is covered.

257

Introduction

Data are collected on Young's modulus and its temperature variation

(0 to 590 K, -460 to 600° F) for forty engineering alloys. Criteria for

including an alloy are: commercial availability, workability and joinability, suitability for low-temperature applications, and sufficient

data to define the temperature dependence of the Young's modulus. The

considered alloys represent four base metals: aluminum, copper, iron,

and nickel

.

The paper's main purpose is to provide a single source of reliable

(± five percent) Young ' s-modul i values for a variety of common alloys

over a wide temperature range. In many cases, Young's moduli of similar

alloys can be deduced from the values reported here.

Data were acquired from three types of sources:

1 . Handbooks

1.1 Cryogenic Materials Data Handbook

1.2 Handbook on Materials for Superconducting Machinery

1 . 3 Metal s Handbook

1.4 Aerospace Structural Metals Handbook

1.5 NERVA Handbook

2. Experimental studies at NBS/Boulder using dynamic methods

3. The scientific-engineering literature

In the second category of sources, dynamic Young ' s-modul us values

between 4 and 300 K have been published for eighteen alloys: Inconel 600^ ,*

Inconel 718^, Inconel X-750^ , nickel steel (5 N i

) ^ , nickel steel (9 N i

) ^

,

*Tradenames are used to describe adequately the considered material. No

endorsement of a particular product by NBS is implied.

259
invar

4

, AISI 304

5

, AISI 31

0

5

, AISI 316

5

, stainless A286

5

, cupernickel 10

6

, 6 7 8 8

cupernickel 30 , aluminum 1100 , aluminum 2014 , aluminum 2219 , aluminum

7 7 7 5083 , aluminum 7005 , and aluminum 7075 . Thus, low-temperature data

from the same laboratory are available in eight of the twelve alloy

groupings used in the study. As described below, reliable high-temperature

values can often be predicted from low-temperature data because of the

linearity of the modul us-temperature curves at higher temperatures.

Description of Young's Modulus

Young's modulus has many pseudonyms: extension modulus, tensile

modulus, tension modulus, elastic modulus, and modulus. Sometimes a

compressive Young's modulus is described. Despite occasional contrary

reports, this is a misnomer. For metals and alloys there is no difference, theoretically or experimental ly, between Young's modulus in extension and in compression.

Young's modulus, E, is usually defined as the ratio of the stress

a

z applied along a rod with axis z to the strain component e

z along the

g same axis. Thus,

Bending of rods is also characterized by Young's modulus. Thus, the

( 1 )

stress along a bent rod with axis z is

10

( 2 )

where x is a coordinate perpendicular to z, and r is the radius of

curvature of the neutral surface near the origin.

260
Velocities of extensional waves in rods are given by

11

v = (E/p

)

1/2

( 3 )

where p is the mass density.

The shear modulus G is related to Young's modulus by the empirical

relationshi

G - 3E/8 (4)

Young's modulus can be correlated with other solid-state parameters. For example,

E = KM0 /d (5)

where K is a constant, M is the average atomic mass, 0 is the Debye temperature, and d is the interatomic nearest-neighbor distance.

Resul ts

Various base metals exhibit a relatively wide range of elastic

moduli, and it is well known that chemical composition is usually the

principal determinant of an alloy's elastic constants. Most other

metallurgical variables have smaller effects on Young's modulus. Thus,

alloys with similar compositions can be grouped. In the present study,

alloys were grouped if their Young's modulus varied less than ± five

percent from the group average over the entire temperature region. In a

few accidental cases, different types of alloys could be grouped within

a ± five percent band. Results of the study are shown in figure 1 and

in table 1. The Monel K curve shows an unusual flatness that probably

indicates inadequate experimental measurements rather than actual material behavior.

261
Modul us-Versus-Temperature Curve Fitting

Almost all data considered in this study were fitted to a the1 4

oretical relationship proposed by Varshni :

E(T) = E(0) - s/[exp (t/T) - 1] (6)

This relationship follows from an Ei nstein-osci 11 ator model of solids.

It contains three adjustable parameters: E(0), the Young's modulus at

zero temperature; t, the Einstein characteristic temperature in the absence of electronic effects; and -s/T, the high-temperature limit of

the temperature derivative dE/dT. T is the temperature in kelvin. Equation (6) was proposed originally to describe the temperature behavior of

1 5 si ngl e-crystal elastic constants of pure metals. It was shown that

it also describes the temperature dependence of the macroscopic elastic

constants, such as Young's modulus of alloys.

Materials described by eq. (6) are called regular , and most materials are. As a class, ferromagnetic materials are irregular; they exhibit modulus-temperature anomalies than can be quite large. The most

irregular material considered in the present study is invar. Because

of magnetostrictive effects, it has a posi ti ve value of dE/dT over a

large temperature region. Several other materials considered show

smaller anomalies in their temperature dependence. These include:

Inconels 600 and X750; austenitic stainless steels 304, 310, 316, and

A286. However, these deviations from regular behavior were less than

the ± five percent bandwidth described above, and they were neglected.

In these materials, most deviations occur at lower temperatures where

magnetic effects are relatively less important.

262
Once parameters in eq. (6) are determined, which requires at least

four good (or three excel 1 ent ) data points, E values can be computed for

the entire temperature range. Because of the approximate linearities of

the E/T curves at higher temperatures , extrapolation of low-temperature

data is quite reliable. The reverse extrapolation is not. In fact,

fitting eq. (6) to data usually requires at least one reliable lowtemperature point below 75 K. Also extrapolation to high temperatures

is less reliable for materials with lower melting points, T . At high

values of T/T , polycrystal 1 i ne materials tend to soften elastically

because of grai n-boundary relaxations. Thus, copper-alloy moduli, which

are extrapolated, may be slightly high.

Once groupings were determined by comparing curves for individual

materials, the group-average curve was determined by refitting eq. (6)

to all the alloys in the group. Each alloy was represented by the same

number of points from its fitted-curve values.

Goodness-of-fi t (average difference between observed value and curve

value) varied considerably among various data sets. Ultrasonic data

were usually excellent, 0.05 percent or less. Static data were usually

poor, several percent or more. In many cases, sets of static data were

discarded because eq. (6) could not be fitted to them. In general, all

readily available data were included unless the data set was highly

scattered, significantly different from the average of all these data

available for the alloy, or showed irregular temperature behavior.

For almost all alloys considered, room-temperature Young's moduli and

dE/dT values are known reliably. Thus, there are criteria for rejecting

poor data sets.

263
Static and Dynamic Values

For most engineering applications, the appropriate Young's moduli

are the static (or zero-frequency) values. These values are obtained

in conventional stress-strain-type experiments associated with mechanicalproperty measurements. Dynamic Young ' s-modul us values are now being

determined more often because less inaccurate values are required, more

physicists are studying mechanical properties of engineering alloys,

and the advantages of dynamic methods are becoming better understood.

These methods, which typically use kHz frequencies for resonance studies

or MHz frequencies for pulse studies, yield dynamic (adiabatic) moduli

that always exceed static (isothermal) moduli. For Young's modulus

1 6

the i nterrelationshi p is

E. = E/(1 -E Ta

2 /C ) (7)

d s'

v s' pv '

where subscripts d and s denote the dynamic and static modulus, T is

temperature, a is linear thermal expansion, and C is specific heat

r per unit volume at constant pressure. Clearly, the difference between

and E

s vanishes at zero temperature for all materials and is small

at all temperatures for 1 ow-thermal -expansion materials such as invar.

In the present collection of Young's moduli, no static-dynamic corrections were applied to the data. The magnitudes of such corrections

are shown for various base metals and various temperatures in table 2.

For the materials and temperatures of concern here this correction is

insignificant, never exceeding one percent.

Observations and Conclusions

(1) Except for invar, all considered alloys have regular temperature

behavior: linear at higher temperatures , decreasing slope with decreasing

temperatures, and relative flatness approaching zero temperature.

264
(2) For the considered temperature region, except for invar, the

modulus-temperature curves are similar and the high-temperature d E/dT

values cluster around -0.06 GPa/K with a la standard deviation of 20 percent.

(3) Aluminum-base alloys have a smaller alloying dependence than

alloys based on nickel, iron, or copper.

(4) Copper-base alloys have the largest alloying dependence, although many copper alloys can be grouped in a ± five percent band.

(5) Only copper-nickel alloys have higher Young's moduli than the

base metal. Alloying usually reduces Young's modulus.

(6) If a magnetic transition occurs, the anomaly size is related

to the magnetic-transition temperature. Invar undergoes a Curie transition near 550 K and shows a large elastic anomaly. The stainless steels

undergo Neel transitions near 50 K and show much smaller elastic anomalies

Acknowl edqment

The study was supported partially by the NBS Office of Standard Refer

ence Data. A. F. Schmidt of NBS encouraged the study and commented on

the manuscript. Both the tabular and graphical data contained in this

paper were contributed by NBS to the American National Standards Committee B31.10 for inclusion in the ANSI B31.10 Code for Cryogenic Piping.

265
References

1. Weston, W. F., Ledbetter, H. M., and Naimon, E. R., Mater. Sci

.

Eng. 20, 185-1 94 (1 975).

2. Weston, W. F. and Ledbetter, H. M., Mater. Sci. Eng. 2_0, 287-290

(1 975).

3. Weston, W. F., Naimon, E. R., and Ledbetter, H. M. , in Properties

of Materials for Liquefied Natural Gas Tankage , ASTM STP 579 , Amer. Soc . Test. Mater., Philadelphia (1970), pp. 397-420.

4. Ledbetter, H. M., Naimon, E. R., and Weston, W. F., in Advances

in Cryogenic Engineering, Vol . 22, Plenum, New York ( 1 977 )

,

pp. 174-181 .

5. Ledbetter, H. M., Weston, W. F., and Naimon, E. R., J. Appl

.

Phys. 46, 3855-3860 (1975).

6. Ledbetter, H. M. and Weston, W. F. , in Ultrasonics Symposium Pro - ceed! ngs , IEEE, New York (1 975), pp. 623-627.

7. Naimon, E. R., Ledbetter, H. M., and Weston, W. F., J. Mater.

Sci. 1_0, 1309-1 31 6 (1 975).

8. Read, D. T. and Ledbetter, H. M., J. Eng. Mater. Technol., forthcomi ng

.

9. Landau, L. D. and Lifshitz, E. M., Theory of Elasticity , Pergamon,

London (1959), p. 13.

10. Ref. 9, p. 73.

11 . Ref. 9, p. 110.

12. Ledbetter, H. M., Mater. Sci. Eng. 2!7, 1 33-136 (1 977).

13. Grimvall , G. and Sjodin, S., Phys. Scr. 1_0, 340-352 (1 974).

14. Varshni, Y. P., Phys. Rev. B 2_, 3952-3958 (1 970).

15. Naimon, E. R., Weston, W. F., and Ledbetter, H. M., Cryogenics 1_4,

246-249 (1974).

16. Ref. 9, p. 17.

266
Table

1.

Young's

modulus

of

groups

of

engineering

alloys

in

units

of

GPa.

Aluminum Alloys Oi cr> oi Oi Oi co co r— ud ix) uo n ro oo .— ooicONrv^Difi<-foc\jr-po^Oico

Copper Alloys ^•^^^r^cocsjcvjf— oc^corvVDLn^t^roNr- oo^coNtoin^ooc\jr- o ,— r— r— r^-r^r— »—OOOOOOOOOOOcncricncr.cr1 cricncncT>a>

Copper O^dCOCOCONtOiOLn^fOCMr-OOOlCONlDlO^COCvj.— o cn CO r\ VO LD cocococooooococococococococococ\jc\jc\jc\jc\jc\jc\jc\jc\jc\jr— i— .— .— .— .

Invar «<frcoc\jc\jcNjcoco«3-Lnvor^cooNi— c\jn«d- ion coOi— NfoimruouDNiOkO fonfonfororonfonronn^^^^^^^L/unini^LnimruniOLnir)

Cu-lONi COCOCOcocococONNtoi^^^tncvj^c^cONtOLO^coc^oo^corvvoinco nr^t<t^-'tt’t^^-’d-’d-’d-^-^^^-ronfOfororororocoo»jcg(\icNjc\jCM

ZZOCO

1 3O

tt'd-^-«d-^'vrrorocdCvJr-oo>cON'Dio^n(\jf— ocicoi^Lfi^-roc\ir-o iovovo^otoi£Koy3vo^CM£)voi/ufunif)ir)LDLni^Lr)i/i^'^«t^«t ,^^'d-^

Monel

K

ooooooooooooooocncno>cocor^r^'X5i-0'^rroc\j l— o co n

COCOCOCOCOOOCOCOCO aDCOCOCOCOCONNNNNrsNNNNNNNND^

Stainless Steel

s

NNNrs^OlDir>r)C\Jr-OiCOlOLfiCOOdOCTirsl/l^C\Jr-Oihsi£)<tfOr-CTiCO OOOOOOOOOOOIOIOIOIOIOIOICOCOCOCOCOCONNNSNS to to C\JC\JC\JC\jCSJC\JC\JC\JC\JC\Jr— r—r— r— r— r— r— »— r- r— «— r— r-r- r-r— r— r— r—

Nickel Steel

s

r^rNrvKrvNtoio^^nr-oaicONtoioncMf-oa>rvtoinntfOf-oo^ ooooooooooooooioioioioicTtoiaia'cococococococoooN

CSJ C\J CSJ CSJ CSJ C\J CSJ C\J C\J CSJ CSJ CSJ CSJ r— r— r— r— r— r— r— r— r— r— r— r— r— r— r— r— r—

Inconel

s Lnini/untn^focxji-ooicotoio^ni— ooicotoio^Nr-ocoNtoion .— .— r— 1— 1— .— .— i— .— r-OOOOOOOOCTiCJ^CnOlOlOiOiCncOCOCOCOCO

CSJCSJCSJCSJCNJCSJCSJCSJCSJCSJCSJCSJCSJC\JCSJCSjCSJCSJr— r-r— r— r— .— r— r—r— r— r— »— i—

Nickel Alloys CNJ C\J C\J C\J CSJ C\J i— 1— OOOiOONtOlO^rOOJ.— OOiCOtOUO^-fOCSJO^CON C\J(VlC\JMM(\JCtJCMC\JCVJr-r-r-i-r-r-f-r-i-r-00 0 00000010>Ol CSJ C\J CSJ C\J CNJ CSJ CSJ CSJ CNJ CNJ CvJ CSJ CNJ CSJ C\J CSJ CSJ Cvj CSJ CSJ CSJ CSJ CNJ CNJ C'J CSl C\J CSJ r— r— r—

Nickel oooooicoNto^foovjooicotDLn^cNjr-OicON^^n.— o co rv to «5j- «d-«d-«cr<3-cocococooorororocsjcsjCNjcsjcsjcsjcsji— t—r—r— •— — .— r— O O O O C\JCSJCSJCSJC\ICSJCSJCSICSJCSJCSJC\ICSJCNJCSJCSJCSJCMCSJCSJCSJCSJCSJCSJCSJC\JC\ICSJC\ICSJCSJ

Temp. (°C) 0-000000 Ooo OOOOOOOOO 00:0 ooooooooo N VO (VJ O CO to CSJ O CO VO CNJ CSMtvOCOO,<M«t\OCOOCSMJ-tOCO O CSJ CVJCSJCSJCSJCSIr— r— r- •*—**»— till r— r— r—r— r— CSJCSICSJCMCSJCOCO

1 1 1 1 1 1 1 1 1 1

267

280

207

199

186

181

171

170

142

126

157

116

300

206

198

185

180

169

168

141

125

156

115

320

204

197

183

179

168

167

140

123

156

114
Table 2. Ratios of dynamic (adiabatic) and static

(isothermal) Young's moduli, E^/E .

Temp. (K) Nickel Iron Copper Alumi num

100 1 .000 1 .000 1 .001 1 .001

200 1 .002 1 .001 1 .001 1 .003

300 1 .003 1 .003 1 .003 1 .005

400 1 .004 1 .004 1 .004 1 .007

500 1 .005 1 .005 1 .006 1 .008

600 1 .006 1 .006 1 .007 1 .010

268
TEMPERATURE.

(spucsnoqj ui ) ejw AlDllSVdd dO SmiTGOW

c<

=3

(—<

O

o

o

o

(spuesnoiji ui) ;s>| AiDliSVTd dO SDiriGOW

269

Figure

1.

Temperature

variation

of

Young's

modulus

for

forty

engineering

alloys.

WELDING AND CASTING

271

EVALUATION OF WELDMENTS AND CASTINGS

FOR LIQUID HELIUM SERVICE

T. A. Whipple and H. I. McHenry

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

In this paper the FY80 results of the NBS program on the evaluation

of weldments and castings for liquid helium service are reviewed. The

welding process study for stainless steel castings and the study on the

toughness of 5183 aluminum weldments were completed during FY80, and

studies on the toughness of stainless steel castings and the fracture

and deformation phenomena in duplex austenite/ferrite structures were

begun. The research planned for FY81 is also reviewed.

273

Introduction

The program to evaluate weldments for structural applications in

superconducting magnet systems [1,2,3] was continued in FY80. Except

for the addition of the evaluation of stainless steel castings for

liquid helium service, the objectives for the program have remained

unchanged from those of FY79:

1. To investigate the metallurgical factors that affect the

toughness of stainless steel weldments and castings at cryogenic

temperatures.

2. To contribute to the development of improved filler metals for

welding stainless steels for liquid helium service.

3. To evaluate the strength and toughness of stainless steel

weldments and castings and aluminum weldments at 4 K.

Research began this year on a study that will take a more fundamental

approach to the first objective, and work has continued on the evaluation

of experimental stainless steel filler metals, evaluation of aluminum

weldments and support of superconducting magnet construction programs.

This report summarizes the results obtained in FY80 and outlines the

research to be conducted in FY81

.

Stainless Steel Weldments

Since previous studies [4,5] have shown that ferrite in weldments

substantially reduces the fracture toughness at cryogenic temperatures

,

substantial work is being conducted on the development of 0% ferrite

filler materials. A number of weldments made with experimental filler

alloys have been tested at 4 K over the past year. The results of

275
these tests are presented later in this report by Elmer, McHenry and

Whipple. Some of these experimental weldments have shown excellent

combinations of strength and fracture toughness. However, some of the

alloys have considerable problems with microfissuring and none are as

yet commercially available.

The study on the effects of the welding process on the cryogenic

properties of 31 6L weldments, which was begun in FY79 [6], was completed

during the past year. The results of 4 K mechanical property testing

are presented later in this report by Whipple and Kotecki. The weldment

made by gas tungsten arc welding (GTAW) was found to have the highest

toughness followed by gas metal arc welding (GMAW) and submerged arc

welding (SAW). This order of weld metal toughness is the same as the

expected order of weld metal cleanliness, indicating that impurities in

the weld metal may have a significant influence on cryogenic properties.

As in past years, a number of qualification and test weldments were

tested for tensile and fracture toughness at 4 K in support of the LCP

magnet design and construction programs. These weldments included a

316L shielded metal arc (SMA) weldment, which was annealed, and 308L

flux core metal arc (FCMA) weldments. The results of these tests are

presented in Tables 1 and 2. In addition, these tables contain data

obtained on a 31 6L electroslag (ES) weldment and a 308 SMA weldment (in

the as-welded and annealed conditions), which were tested under separate

contracts with the MFTF and the ANL-MHD projects, respectively. These

data are included here because the weldments are being used as part of

the overall study of 4 K fracture and deformation of stainless steel

weldments at 4 K, which is included in this program.

276
As mentioned previously, a major fundamental study of fracture and

deformation in duplex austenite/ferri te stainless steels at cryogenic

temperatures has been undertaken. The objective of this investigation

is to relate the structure of these materials to the cryogenic properties.

During FY80 two initial parts of this investigation were completed and

are presented later in this report. The first, by Brown, presents a

detailed literature review, critical analysis and observations concerning

structure development in duplex austenite/ ferrite stainless steel weldments

and castings. This study is necessary to define structural parameters

to be used in deformation modeling and to understand effects that may

arise due to such factors as microsegregation. This knowledge will also

be necessary for the eventual modification of structural parameters to

optimize mechanical properties. The second paper, by Whipple and Brown,

presents the results of a metal lographic investigation of deformation

and fracture mechanisms in duplex austenite/ ferrite stainless steel

weldments and castings. In this study, the role of ferrite in 4 K

deformation and fracture is examined with particular emphasis on possible

fracture modes. The fracture-mode study has pointed out the advantages

to be gained in studying castings as a model material that will give

indications of the behavior of weldments. Because castings have a much

coarser ferrite structure than weldments, their deformation and fracture

phenomena are much more easily observed.

Stainless Steel Castings

The increased interest in the use of stainless steel castings for

structural components of superconducting magnets has resulted in a

program to evaluate castings for cryogenic service. Ferrite content and

277
nitrogen concentration have been found to have a significant influence

on the cryogenic properties of stainless steel weldments [4,5] and it

seems reasonable to expect that the same will be true of castings because

of the similarities in their microstructures. Therefore, the program

has been designed to test both of these effects. Two series of CF8M

castings have been produced, one with a constant nitrogen content and

varied ferrite content and one with a constant ferrite content and a

varied nitrogen content. The production and initial testing of these

castings is presented later in this report by Finch.

In FY80, 4 K mechanical property tests were conducted on some CF8M

and CF8 castings under separate contracts. The data from these tests

are presented in Tables 3 and 4. These castings have been used as part

of the deformation and fracture mechanism study being conducted in this

program, the results of which are presented later in this report by

Whipple and Brown.

Aluminum Weldments

The program on the evaluation of 5183 welds in 5083 aluminum, which

was begun in FY79 [7], was completed this year with 4 K fracture toughness

testing. It was found that the toughness of aluminum weldments is not

sensitive to welding procedure or supplier. The results of this program

are presented later in this report by Elmer and McHenry.

Test-Method Development

In FY79 work had begun on the evaluation of a simple test method

for estimating Jj

c [8]. The aim of that work was to develop a simple

and inexpensive test method for weld qualification for 4 K service. The

initial results indicated that the test worked well for the base metals

tested. In the past year the test method was evaluated for use on weldments. The weldments tested were those from the welding process study,

278
mentioned previously. The results are shown graphically in figure 1,

which is a plot of the Jj estimated by this procedure versus the Jj

c measured by the single-specimen J-integral procedure. From this figure

it is evident that, in the case of weldments, the correlation between

the estimated Jj

c and the actual Jj

c is very poor. On the basis of

these results, it was determined that the test method is not useful for

weld property evaluation. This is probably due to the large grain size

of weldments compared with the small ligament of the specimens used for

this test.

Future Work

The major portion of the research effort in FV81 will be devoted to

the previously mentioned deformation and fracture study on duplex austenite/

ferrite stainless steel weldments and castings at 4 K. The objective of

this work is to gain a much more comprehensive understanding of factors

influencing the mechanical properties, including the causes of the

observed large variability in weld metal properties. Specifically, this

research will include the following efforts:

1. A theoretical model of duplex austenite/ferrite stainless

steel deformation at 4 K will be developed from a two-phase

deformation modeling approach. This model is intended to

define quantitatively the effects of ferrite content, ferrite

morphology, and phase properties on deformation.

2. The CF8M castings produced in FY80 will be used to verify and

refine the above model and to make additional fracture and

deformation mechanism observations.

279
3. Weldments of 31 6L stainless steel weldments will be made and

tested with different welding processes and ferrite contents

to extend the model to the finer structure in weldments.

4. Extensive metallographic investigations will be conducted in

support of the above program. Optical microscopy will be used

to define and measure microstructural parameters. Transmission

electron microscopy will be used extensively to investigate

factors such as orientation relationships, interface structures

and dislocation substructures or deformed and undeformed

specimens. Scanning electron microscopy will continue to be

used to correlate microstructure to fracture and deformation

micromechani sms

.

5. Microstrain tensile testing, using high-sensitivity strain

measuring equipment, will be conducted to establish microdeformation

phenomena further.

In addition to and as part of the above program, a COD test methodology

will be developed. This test method will use three-point bend specimens

and will be capable of measuring Jj

c> It is expected that this form of

testing will significantly reduce the cost of 4 K fracture toughness

testing. The 4 K mechanical property testing in support of superconducting

magnet design and construction programs will continue in FY81

.

280
REFERENCES

1. H. I. McHenry, "Evaluation of Stainless Steel Weld Metals at Cryogenic

Temperatures," in Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures-I , NBSIR 78-884 (1978), pp. 159-166.

2. H. I. McHenry, D. T. Read, and P. A. Steinmeyer, "Evaluation of

Stainless Steel Weld Metals at Cryogenic Temperatures , " in Material

s

Studies for Magnetic Fusion Energy Applications at Low Temperatures

IJ_, NBSIR 79-1 609 (1 979), pp. 299-308.

3. H. I. McHenry and T. A. Whipple, "Weldments for Liquid Helium

Service," in Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627 (1980), pp. 155-165.

4. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr.,

"Metallurgical Factors Affecting the Toughness of 316L SMA Weldments

at Cryogenic Temperatures," Material Studies for Magnetic Fusion

Energy Applications at Low Temperatures-I

I

, NBSIR 79-1609 (1979)

pp. 313-352.

5. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic Stainless Steel Weld Metals: Part I - Role of Ferrite,"

Weld. J., Weld. Res. Suppl

.

57, (1978) pp. 325-s-333-s.

6. D. J. Kotecki, "Weld Process Study for 316L Stainless Steel Weld

Metal at Cryogenic Temperatures," in Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures-III , NBSIR 80-1627

(1 980), pp. 197-204.

7. R. A. Kelsey and L. N. Mueller, "Tensile and Notch-Tensile Properties

at Room Temperature and 4K of Welds in Aluminum Alloy 5083-0

Plate," in Material Studies for Magnetic Fusion Energy Applications

at Low Temperatures-III , NBSIR 80-1627, (1980) pp. 217-244.

281
8. T. A. Whipple and H. I. McHenry, "Note on a Simple Test Method for

Estimating Jj

c

>" i n Materials Studies for Magnetic Fusion Energy

Applications at Low Temperatures-1 1

1

, NBSIR 80-1627 (1980), pp. 245-256.

282
Table 1. Tensile properties of qualification

weldments tested at 4 K.

Wei d Ferri te

,

Yield Ul timate Elongation, Reduction

Type % Strength

,

MPa (ksi)

Strength,

MPa (ksi)

% in Area,

%

31 6L SMA 0 636 939 13 11 685 1349 40 24

650 919 13 20

727 1315 46 43

674 (97.8) 1130 (164.0) 28 25

31 6L 4.5 700 1436 34 23

El ectrosl ag 680 1294 38 25

690 (100.0) 1365 (197.4) 36 24

308L 3 483 1168 22 16

FCMA 431 1395 29 19

0.60N 460 1374 32 21

458 (66.4) 1312 (190.3) 28 19

308L 3 541 1418 31 22

FCMA 574 1219 21 16

0.8N 558 (80.9) 1319 (191 .3) 26 19
Table 2. Fracture properties of qualification

weldments at 4 K.

Weld Ferri te,

of

<_> p

—

r

"D K Ic

(J), Fatigue Crack Growth

/o kj/m^ o MPav/ m—

C n

2 (in-lb/in

2

)

( ksi /in’)

metri

(Engl i sh

units

units)'

31 6L SMA 0 241

(annealed) 288

111

267 (1526) 247 (224) 0.501 x 10"? 2.83

(0.263 x 10 )

31 6L 4.5 174

El ectrosl ag 140

157 ( 897) 189 (1 72)

308L 3 80.3

q FCMA 77 0.489 x 10'

n (0.281 x 10

U

)

3.69

0. 8N 80.3

79.2 ( 453) 1 34 (1 22)

308 SMA 12 46.0 0.254 x 1 0~

g 3.84

32.2 (0.163 x 10" ) 3.88

39.1 ( 224) 94.4 (85.9)

308 SMA 0 328 0.196 x 1 0~

3.78

Anneal ed 350 (0.846 x 10' ) 3.34

339 (1 938) 278 (253)

1. Constants fit the equation: da/dN = CAK

n

2. Metric units: AK in MPav'mT da/dN in mm/cycle

English units: AK in ksi/Tn, da/dN in in/cycle

284
Table 3. Tensile properties of stainless

steel castings at 4 K.

Material Ferrite,

1o

Yi el d

Strength

,

MPa (ksi)

U1 timate

Strength

,

MPa (ksi)

El ongation,

%

CF8 0 413 (60) 125 (181) 46.1

8 689 (100) 920 (133) 20.2

14.5 470 (68) 1512 (219) 47.6

12 578 (84) 753 (109) 10.7

CF8M 9.0 621 (90) 1317 (191) 41.0

469 (68) 1338 (194) 44.0

285
Table 4. Fracture properties of stainless-steel

castings at 4 K.

Material Ferrite, J Ic’ 0 K tr (j), Fatigue Crack Growth

% kJ/m

2 ? MPa/rrP C 2

n (in-lb/in ) (ksi/in) metric units

CF8 0 481 (2750) 331 (301)

8 175 (1000) 199 081) -- -- 14.5 69.3 (396) 125 (114) -- -- 12 31.5 (180) 84.5 (76.9) -- --

CF8M 9.0 302

403

280

-12

2.35 x 10 5.21

372

339 (1938) 278 T253J

1. Constants fit the equation: da/dN = CAK

n

2. Metric units; AK in MPa/mV da/dN in mm/cycle

286
500 CO CM

O

o

m

o

©

©

©

CO

o

©

CM

O

o

CM

O

S- I—

05 CO

Q.£ S- O O

CJ> M—

cus-3

CD

Ll_

jUi/r* ‘aiansvaw °'r

287

ESTIMATED,

kJ/m

STRENGTH AND TOUGHNESS OF FULLY AUSTENITIC STAINLESS STEEL

FILLER METALS AT 4 K

J. W. Elmer, H. I. McHenry, and T. A. Whipple

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Several experimental fully austenitic stainless steel welds have

been prepared by filler metal suppliers in the United States, Austria,

Sweden and the Soviet Union. The strength and toughness of these welds

have been tested and are reported at 4 K. Some Charpy impact toughness

and tensile data at temperatures other than 4 K are presented. These

data were reported by the suppliers. Four of the eight welds that were

tested had strength-toughness combinations comparable to that of austenitic base metals. Of these four welds, three did not exhibit microfissuring that is generally present in fully austenitic welds.

289

INTRODUCTION

Stainless steel weldments are used for the main structural members

of the superconducting magnet systems for fusion energy devices. Annealed austenitic stainless steels of the A I S I 300 series are currently

being used in weldments with thicknesses up to 125 mm. The weldments

operate at high stress levels in liquid helium at 4 K. The weld metal

that has been most thoroughly evaluated and is most frequently used for

liquid helium service is AWS 31 6L deposited by the shielded metal -arc

(SMA) process. Fracture toughness tests on 31 6L welds at 4 K indicate

that toughness decreases with increasing ferrite content^ and with

2 increasing grain size in ferrite-free welds. In general, the toughness

3

of 316L welds varies significantly for the different welding processes,

and the best assurance of satisfactory toughness is obtained with a

ferrite-free weld deposited with a low-heat input. Problems arise

4 5 because ferrite-free welds are susceptible to microfissuring ’ and

low-heat inputs result in low deposition rates.

The purpose of this program is to conduct a preliminary evaluation

of the strength and toughness at 4 K of several proprietary fully austenitic weld metals provided in the form of as-welded test plates by

filler metal suppliers in the United States, Austria, Sweden, and the

Soviet Union. The evaluation is limited to the specific welding process

and procedures chosen by the suppliers to prepare the test plates.

Promising grades will be identified, but further evaluation will be

needed to assess their suitability for structural members of superconducting magnets.

291
MATERIALS

Test weldments were provided for evaluation by five filler metal

suppliers. The chemistry of the weld metals is summarized in Table I.

The welds are identified by their chemistry: the numbers indicate the

Cr-Ni-Mn content and any additional characteristic alloying elements

follow, e.g., alloy 18-20-6 Nb is 18% Cr, 20% Ni, 6% Mn, and contains Nb.

The welding process and procedures used to prepare the test weldments

were selected by the suppliers. The available information is summarized

in Table II. Tensile and Charpy impact data provided by suppliers

4 and 5 are summarized in Table III. Microfi ssuring was observed on the

fracture surface of all the welds except for the 20-16-7 Mo-W submerged

arc weld (SAW), 17-16-4 Mo (SMA), and 18-20-6 Nb gas metal-arc (GMA)

wel ds

.

PROCEDURES

Tensile Testing

Tensile testing was conducted on al 1 -wel d-metal specimens by means

of a screw-driven testing machine equipped with a cryostat to perform

the tests in liquid helium. A 0.635-cm (0.250-in)-diameter tensile bar

was used, and the strain was measured over a 25-mm-gage length with a

clip-gage extensometer mounted on the tensile specimen. A constant

crosshead displacement rate of 0.008 mm/s was used for all specimens.

Fracture Toughness Testing

All fracture toughness testing was performed using the computerg aided, single-specimen, J-integral test facility described by Read.

The testing was done in liquid He at a temperature of 4.0 K using a

standard 2.54-cm ( 1 . 00-in )-thick ASTM E399 compact tension specimen.

The specimen was removed with the plane of the notch perpendicul ar to

292
the surface of the plate and centered along the fusion zone. In some

instances, the weldment was not thick enough to remove a 2 . 54-cm-thick

compact tension specimen. In these cases, the planar dimensions were

kept the same as for a 2. 54-cm-thick specimen and only the thickness was

reduced. The specimens were precracked in tension-tension fatigue to a

depth of 3.05 cm (1.20 in) prior to testing. Displacements were measured

with a clip-gage extensometer mounted at the load line.

RESULTS AND DISCUSSION

The tensile properties of the fully austenitic welds at 4 K are

summarized in Table IV. The yield strength ranged from 606 to 1110 MPa.

Much of the variation in yield strength can be attributed to variations

in the C-and-N contents. These interstitials are effective strengtheners

at low temperatures . However, the correlation of yield strength with

interstitial content reported by Tobler and Reed'

7 for AISI 304 base

metal was not observed. Apparently, other factors such as grain size,

oxygen content, and overall chemistry are important contributors to

strength. For example, four welds with the same carbon plus nitrogen

(0.08%) had yield strengths ranging from 606 to 814 MPa. Similarly, the

weld ductility, as measured by reduction of area, generally decreased

with increasing C plus N. For example, the three welds with less than

20% reduction of area had the highest C-plus-N contents.

The fracture toughness values of the fully austenitic welds at 4 K

are summarized in Table V. As a group, these welds had high fracture

toughness relative to previous results on AWS 31 6L welds with and with1 2 out ferrite. ’ A comparison of the present results with the trend line

g of Read and Reed is shown in Figure 1, a plot of fracture toughness

versus yield strength. Note that the strength and toughness combinations

293
for 20-16-7 Mo-W (SAW), 1 7-1 6-4 Mo (SMA), 18-20-6 Nb (GMA) and 16-16-4 Mo

(FCAW) each are close to the base-metal trend line. This is noteworthy

because four different welding processes are used, a wide range of

strengths is covered, and three of the alloys (all except 16-16-4 Mo)

did not exhibit microfissures . The similarity in chemistry of these

alloys suggests that a fully austenitic alloy with 16 to 20% each of Cr

and Ni, 4 to 7% Mn and refractory elements may provide the properties

needed for liquid helium service.

SUMMARY AND CONCLUSION

A preliminary evaluation of the strength and toughness of several

proprietary fully austenitic weld metals has been conducted. The results for 20-16-7 Mo-W, 17-16-4 Mo, 18-20-6 Nb and 16-16-4 Mo indicate

that strength and toughness combinations comparable to the properties of

austenitic stainless steel base metals are possible in fully austenitic

weld deposits. Further investigation is needed to identify the optimum

chemistry and welding process for achieving microfissure-free welds that

have strength and toughness properties needed for structural applications in superconducting magnets.

294
REFERENCES

1. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas, Jr.,

"Metallurgical Factors Affecting the Toughness of 31 6L SMA Weldments at Cryogenic Temperatures," Weld. J. Res. Suppl . , Vol . 59,

No. 4, pp. 1 04s- 113s (1980).

2. T. A. Whipple, H. I. McHenry, and D. T. Read, "Fracture Behavior of

Ferrite Free Stainless Steel Welds in Liquid Helium," Material

s

Studies for Magnetic Fusion Energy Applications at Low Temper -

atures-III , NBSIR 80-1627, National Bureau of Standards, Boulder,

Colorado, pp. 167-194 (1980).

3. T. A. Whipple and D. J. Kotecki , "Weld Process Study for 316L

Stainless Steel Weld Metal for Liquid Helium Service," this volume,

pp. 303-321 (1981).

4. C. D. Lundin, W. T. Delong, and D. F. Spond, "Ferrite-Fissuring

Relationship in Austenitic Stainless Steel Weld Metals," Weld. J.

Res. Suppl., Vol. 54, No. 8, pp. 241 s -246s (1975).

5. C. D. Lundin and D. F. Spond, "The Nature and Morphology of Fissures in Austenitic Stainless Steel Weld Metals," Weld. J., Vol. 55,

No. 11, pp. 3565-3675 (1976).

6. D. T. Read, "The Computer-Aided J-Integral Test Facility at NBS,"

Materials Studies For Magnetic Fusion Energy Applications at Low

Temperatures-III , NBSIR 80-1627, National Bureau of Standards,

Boulder, Colorado pp. 205-216 (1980).

7. R. L. Tobler and R. P. Reed, "Interstitial Carbon and Nitrogen

Effects on the Tensile and Fracture Parameters of AISI 304 Stainless Steels," Materials Studies for Magnetic Fusion Energy Applica -

tions at Low Temperature-III , NBSIR 80-1627, National Bureau of

Standards, Boulder, Colorado pp. 17-48 (1980).

295
8. D. T. Read and R. P. Reed, "Fracture and Strength Properties of

Selected Austenitic Stainless Steels at Cryogenic Temperatures ,

Materials Studies for Magnetic Fusion Energy Applications at

Low Temperature-II , NBSIR 79-1 609, National Bureau of Standards,

Boulder, Colorado pp. 79-122 (1979).

296
Table

I.

Chemical

Analysis

of

Test

Welds*

297

FCMA

=

Flux

core

metal

arc
Table II. Welding Parameters

Suppl ier Process Passes T hi ckness

,

mm Current

,

A

Vol tage

,

V Travel Speed,mm/s

1 SAW 2 19 — — —

2 SMA 18 25 130 22 —

3 GMA — 13 220 30 5

4 GMA 20 25 150 — 0.5

5 FCMA — — — — —

5 SMA — — — — —

298
Table

III.

Mechanical

Property

Data

Provided

By

the

Suppliers

of

the

Test

Weldments

M-OEO "

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+-> O)

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4— CO CO CO CO ON CO CO ON 1 1 CO CO CO

UJ 5-5

O 1 1 O „ n —s LO . cr> ,—. -E •!— • LO • a»

+-> CO C\J CO cr» CD

03 -X LO 1

<d e v V **. ^ E

<u O) E >

+-> cC o

— T\ E

"O rd (

cr> CO

O i— D- CO CNJ CO • r— a) 2: cr> c

dJ CO

•i— E E

>- -E fd CO CL

=3

X

O LU 1 1 1 1 1 1 1 1 o o CO CO CNJ 1— E 1 1 1 1 1 1 1 1 LO LO CNJ O

>— E 1 1 1 1 1 1 1 1 • • • • •

r* 4-> fd 1— f— r— i

r— r—

J

^ X

—X—

U E

-t-> «-x LO • rd QJ OVi- . CO • Q. +-> e co i

E fd <D -X or i

r-^ i—

_l e —+-> — ‘ ' >, go CLE

<u rd •—

1 JC -Q ^ —

/" * X

xXX

—>, ^—^—

1

c_> ,co .— OJ CTi r— CO LO LO LO •=d- CNJ 1

CO ro iCO CO coO ON LO o cr> LO LO LO LO E D- LO CNJ cr> >1 +-> V s |— V ^ N f— V ^ ^ V v ^ V - *— •X <u LO CNJ 03 <4— x - V - i—

dJE

LlI on r**^ 1 «• CO LO CO CO CO co LO LO o r\ •— i— CO C\J ON CO CO O O LO co CO 00 r^x 0) T> r— r— f— |— ( i r— i—— E

Z3

+-> rd 1— 1— LD 1— ld rs I— LD

N h- r^ CNJ LO r^x E Cd cd cd on a OIN cd on r-x cd i— r-x dJ r— r— r— i— r— i—

CLE

at h-

•r— o o I— _Q

cr» LO LO LO ON i I | | 1 1 i <c

>3 LO lo o o o LO LO o O i— r— OJ CNJ CNJ i— i— Ll. 

1 1 < 1 1 1 1 | i

00 LO CJ 00 CO CO CO LO c i— Ll_ r— E

CU

•i

i—

Ql

a.

zs GO LO CO CO CO LO 299
Table

IV.

Tensile

Properties

of

Fully

Austenitic

Welds

at4

O "

•I- (O

+-> cu o sC3 <c -a

a> M—

cc o

o E

+-> LD

ra c\j cn

COco O cn co cn co

OO I— CM I— CM CM

co

«=J" CD CD cn I"- O CO

CO CM CM CM OO CM

„—

„

.c: ' ^—

.

^

^ -s LD ^—-x CP) — +-> »r- | 00 1

cn • lo 0 CD co LD 1—

*

o CM cn LO o

i

11

1

cn i

CO 1

ai —

&_

4-> oo

"O

1

O ^d- CD cn co LD LD

CU ro i

(

cn co 00 O i

•r- Q_ > S

1 CO CD CO CD o LD

+-> • rCT) CO

aj —

S4-> OO

aj

CO CM 00 LD cn ,

LO C 03 C\J CD •^d" CO r-^ OO CO

<D Q_ CM O cn 1— oo

i— s: 1— 1— 1— 1— 1— •— 1—

300
Table V: Fracture Toughness of Fully Austenitic Welds at 4 K.

Suppl i er Alloy Process

J/m

2

J lc’

( i n* 1 b/i

n

2

) MPa/m

k t > Ic

(ksi/fn)

1 20-16-7 Mo-N SAW 116 (660) 163 (148)

77 (440) 132 (120)

AVERAGE 97 S' LO LO 148 troy

2 17-16-4 Mo SMA 240 (1370) 234 (213)

306 (1748 264 (240)

AVERAGE 273 (1 560) 250 (227)

3 18-20-6 Nb GMA 130 (740) 171 (156)

AVERAGE 350 (2000) 282 (257)

240 (1370) 227 T207T

3 18-20-6 Ti GMA 362 (2070) 287 (267)

4 18-16-9 GMA 65 (370) 121 (110)

5 16-16-4 Mo FCMA 238 (1 362) 239 (217)

5 16-15-4 Mo GMA 114 (650) 170 (155)

301
UJ/^BdW ‘(D

0|

>l ‘SSBNHOnOl BUniOVUd

302

Figure

1.

4

K

fracture

toughness

versus

yield

strength

for

fully

austenitic

welds.

Data

are

compared

with

the

4

K

austenitic

base

metal

fracture

toughness

trend

from

a

separate

investigation.
WELD PROCESS STUDY FOR 31 6L STAINLESS STEEL

WELD METAL FOR LIQUID HELIUM SERVICE

T. A. Whipple

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

D. J. Kotecki

Teledyne McKay

York, Pennsylvania

ABSTRACT

This study was conducted to determine the effects of welding process

choice on the cryogenic properties of 316L stainless steel welds. Six

weldments were produced using two welding wires and three welding processes.

The weldments were impact tested down to 77 K and tensile and fracture

toughness tested down to 4 K. The best properties obtained were from a

GTA weld, followed by GMA welds; SA welds had the poorest properties.

This variation in properties was attributed to the cleanliness of the

weld metal, which is dependent on the welding process and parameters

sel ected.

303

INTRODUCTION

To date, most structures intended for liquid helium service have

been welded with the shielded metal arc (SMA) welding process. This is

mainly due to the lack of cryogenic mechanical property data on stainless

steel weldments using higher deposition rate processes. The purpose of

this program was to assess the effect of welding process choice and, to

a lesser extent, the effect of welding parameter choice on the cryogenic

toughness of type 31 6L weld metal at two nominal ferrite 1 evel s--approximately

5 ferrite number (FN) and 0 FN. This was intended to add to the data

base on different welding processes and to assess the factors which

cause differences between welding processes.

MATERIALS

Type 316L bare wire is fairly easy to obtain with ferrite calculated

by the DeLong Diagram [1] at about 5 FN . But unless high-nitrogen

content is acceptable, type 31 6L bare wire calculated to have 0 FN is

rare. From previous work [2,3], it is apparent that high-nitrogen

content has a detrimental effect on cryogenic toughness. In order to

begin with a low nitrogen wire of calculated 0 FN , a tubular wire of

special composition was produced for this work.

For the weldments intended to contain 5 FN, a heat of 0.24 cm

(3/32 in) diameter 31 6L stainless steel bare wire with the following

composition was selected:

C Mn Si Cr Ni Mo _P _S _N Co Cu

0.021 1.58 0.50 18.61 1 2.79 2.1 7 0.01 9 0.002 0.033 0.08 0.1 1

305
The ferrite number calculated from the mill analysis using the DeLong

Diagram is 5.2 FN.

A tubular wire was produced deliberately aiming to exceed the 11.0

to 14.0 percent nickel range for type 3]6L wire in AWS A 5.9, so that

0 FN might be obtained. The aim composition of the wire was:

C Mn Si Cr Ni Mo JN

0.02 2.3 0.4 18.5 15.5 2.3 0.02

The calculated ferrite by the DeLong Diagram for this composition

is 0.2 FN. Samples were taken from four locations within this lot. Gas

tungsten arc (GTA) melts were prepared and checked for nickel, chromium,

and molybdenum homogeneity, which proved to be sati sfactory

.

PROCEDURES

Welding

Six weldments were produced for cryogenic testing. All welds were

made in 2.54 cm (1.0 in) thick 304L stainless steel plate in the flat

position. The basic joint preparation was a 90° included angle double

vee which was symmetrical about the plate midthickness. Three weldments

were made with each of the welding wires described above. The bare wire

was used for a GTA weldment, a gas metal arc (GMA) weldment, and a

submerged arc (SA) weldment. The tubular wire was used for a GMA weldment

and two SA weldments, one with normal heat input and one with high-heat

input. Details of the welding procedures, ferrite measurements, and

fissure bend testing have been reported previously [4].

306
Mechanical Testing

The weldments were evaluated using tensile tests, Charpy V-Notch

(CVN) impact tests and fracture toughness tests. The tensile tests were

conducted at 295 K and 4 K; CVN tests were conducted at 77 K, 116 K,

144 K, and 172 K, and fracture toughness tests were done at 4 K.

Tensile tests were conducted at 4 K using 2.5 cm (1 in) gage

length and 0.64 cm (0.25 in) diameter specimens. The specimens were

oriented along the length of the weldment so that the entire sample was

welded metal. The test procedures and the cryostat for the low temperature

tests have been described by Reed [5]. The 295 K tensile tests were

done on standard 1.28 cm (0.505 in) diameter all weld metal specimens.

CVN impact tests were conducted on 12 specimens from each weldment.

The specimens were machined such that the notch was perpendicular to the

plate surface and the crack propagated along the weld centerline.

Fracture toughness tests were conducted at 4 K using the ASTM E399

compact specimen with a 5.1 cm (2 in) width and a 2.5 cm (1 in) thickness.

The specimens were oriented so that loading was perpendicular to the

weldment and crack extension was along the weldment centerline. The

single-specimen J-integral procedure [6] was used to determine Jj , the

value of J at the onset of crack extension. The plane strain fracture

toughness, Kj

c (J), was determined from Jj using the relationship

proposed by Begley and Landes [7]:

where E is Young's modulus and v is Poisson's ratio. Tests were conducted

in an el ectrohydraul ic fatigue machine equipped with the cryostat described

by Fowlkes and Tobler [8],

307
All tests were conducted using the computer aided J-integral testing

facility described by Read [9].

RESULTS AND DISCUSSION

Chemical composition and ferrite measurements were made on all weld

metal deposits. The results of these measurements are given in table 1.

Note that the SA weldments generally have higher concentrations of

residual elements; of particular interest are silicon and oxygen. The

higher concentration of these elements results from the flux used in SA

welding. Also, note that the ferrite numbers do not correlate well with

either the ferrite numbers calculated from the all -wel d-metal deposit

composition or the ferrite numbers calculated from the chemical composition

of the welding wire. As a result of this unpredictabil ity in ferrite

contents, the original objective of obtaining three weldments with 5 FN

and three with 0 FN was not achieved.

Ferrite measurements made on the actual double-vee weldments, which

were tensile and fracture toughness tested at 4 K, are shown in table 2.

These measurements were made with a ferrite meter that samples only a

small area for ferrite content. (Note: These measurements were made in

percent ferrite, which at these ferrite levels is approximately equal to

FN.) By using this instrument, it is possible to see the variation in

ferrite content through the thickness of the weldment. These results

show that, although the average ferrite content of the double-vee weldment

is very close to that of the all weld-metal pad, there is a significant

variation in ferrite content through the thickness of the weldment.

Generally, higher ferrite contents are found in the center of the

weldment, as would be expected from the higher degree of base metal

308
dilution in this area. However, two of the weldments deviated from this

expected trend. The 31 6L bare wire GMA weldment had higher ferrite on

one side of the double-vee than on the other side. The 316L tubular

wire SA weldment with high heat input had higher ferrite near the plate

surfaces than in the center of the weldment. In the case of the high

heat input SA weldment this effect can be explained by the fact that

this weldment had very large passes, which resulted in considerable

remelting of the previous pass and good mixing of the fused base metal.

As a result of the primary austenite solidification mode, ferrite stabilizing

elements were segregated to the top of the weldment. No explanation is

offered for the curious result seen in the GMA weldment.

The above results on ferrite distribution were qualitatively confirmed

by metallographic examination. Figures la and lb show micrographs from

locations C and E, respectively (see table 2). This sample was prepared

using an electrolytic 10 N KOH etch, which selectively deposits a film

on the ferrite phase. The weldments were also examined after color

etching with a solution of 125 ml H

2 0, 50 ml HC1 , and 2 g Na^SO^, which

was used to show whether the solidification mode was primary austenite

or primary ferrite. The solidification mode is generally considered to

be important with regard to microfissuring in the weldment [10] and it

may also have an important effect on cryogenic toughness because it

affects the continuity of the ferrite phase [11]. In general, the examination with the color etch showed that the areas that contained

greater than two percent ferrite solidified as primary ferrite and those

areas with less than two percent ferrite solidified as primary austenite.

The results of CVN impact tests are shown in table 3. From the

very low lateral expansion it is evident that the two normal heat input

309
SA weldments have a low toughness. Although CVN impact results at 77 K

are not a good indication of 4 K fracture toughness [12], it was decided

that the lateral expansions of these two SA weldments were so poor that

4 K testing was unwarranted.

The 295 K and 4 K tensile properties of all weld metal specimens

are shown in table 4. The 31 6L bare wire weldments had higher 4 K yield

and ultimate strengths than the tubular wire weldments. This difference

was large at 4 K and was probably due to the extremely low nitrogen

content of the tubular wire weldments (see table 1).

The results of 4 K fracture toughness tests are shown in table 5.

The highest toughness was achieved with the 316L bare wire 6TA weldment.

The two GMA weldments had almost the same fracture toughness even though

the bare wire GMA weldment had a significantly higher strength. The

tubular wire high heat input SA weldment had poor fracture toughness as

well as a low yield strength.

As noted previously, the 4 K fracture toughness of stainless steel

weldments has been correlated with both ferrite content and nitrogen

concentration. The weldments tested in this study do not have significantly different ferrite contents, and the weldments with lower nitrogen

concentrations also had lower fracture toughness, which is against the

expected trend. The major difference in the weldments is the welding

process used and it seems likely that this is the cause of the wide

variation in fracture toughness.

Of the three processes used, GTA, GMA, and SA, it would be expected

that the GTA process would produce a "clean" weld and the SA process

would result in a "dirty" weld. In this case, clean and dirty refer to

the amount of nonmetal lie inclusions in the weld metal. The SA welding

310
process uses a flux that should result in a higher inclusion count than

processes that use gas shielding. Micrographs on unetched samples from

the bare wire GTA and the tubular wire SA weldment are shown in figures

2a and 2b, respectively. This figure qualitatively demonstrates that

the SA weldment had a much higher inclusion count. It is well known

that a high concentration of hard second phase particles can significantly

reduce fracture toughness, and this appears to be the cause of the

observed variation in fracture toughness of the welds tested in this

study.

SUMMARY

Six 316L stainless steel weldments, which were produced by three

welding processes with two welding wires, were evaluated for cryogenic

mechanical properties. The best combination of 4 K yield strength and

fracture toughness was achieved with the GTA process using 31 6L bare

wire. This weldment had both the highest strength and highest fracture

toughness of the welds tested. SA welding was found to result in both

low strength and low toughness, attributed to the significantly "dirtier"

weld produced by this process.

311
REFERENCES

1. W. T. DeLong, "Ferrite in Austenitic Stainless Steel Weld Metal,"

Weld. J., Res. Suppl . 53, 273s-275s (1 974).

2. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic Stainless Steel Weld Metals: Part I: Role of Ferrite,"

Weld. J., Res. Suppl. 58, 325s-333s (1978).

3. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic stainless Steel Weld Metals: Part II: Role of Nitrogen,"

Weld. J., Res. Suppl. J58, 34s-44s (1 979).

4. D. J. Kotecki, "Weld Process Study for 316L Stainless Steel Weld

Metals at Cryogenic Temperatures," in Materials Studies for Magnetic

Fusion Energy Appl ications at Low Tempera tures-1 1

1

, NBSIR 80-1 627,

National Bureau of Standards, Boulder, Colorado, pp. 195-204 (1980).

5. R. P. Reed, "A Cryostat for Tensile Tests in the Temperature Range

300 to 4 K, " in Advances i n Cryogenic Engineering , Vol . 7 , PI enum

Press, New York pp. 448-454 ( 1 962)

.

6. G. A. Clarke, W. R. Andrews, P. C. Paris, and D. W. Schmidt, in

Mechanics of Crack Growth , Proceedings of 1974 National Symposium

on Fracture Mechanics, ASTM STP 590, American Society for Testing

and Materials, Philadelphia, Pennsylvania, pp. 27-42 (1976).

7. J. A. Begley and J. D. Landes, in Fracture Toughness , Proceedings

of 1971 National Symposium on Fracture Mechanics, Part II, ASTM STP

51, American Society for Testing and Materials, Philadelphia,

Pennsylvania, pp. 1-20 (1972).

8. C. W. Fowlkes and R. L. Tobler, "Fracture Testing and Results for a Ti 6A1-4V Alloy at Liquid Helium Temperatures , " Eng. Fract. Mech.,

8, 487-500 (1976).

9. D. T. Read, "The Computer-aided J-integral Test Facility at NBS,"

i n Material s Studi es for Magnetic Fusion Energy Appl ications at Low

Temperatures-1 1

1

, NBSIR 80-1 627, National Bureau of Standards,

Boulder Colorado, pp. 205-21 6 (1 980).

10. T. Takalo, N. Suntala, and T. Moisio, "Austenitic Solidification

Mode in Austenitic Stainless Steel Welds," Met. Trans. A, 10A,

pp. 1173-1181 (1979).

11. T. A. Whipple and E. L. Brown, "Deformation and Fracture of Stainless

Steel Weldments and Castings at Low Temperatures , " in Material

s

Studi es for Magnetic Fusion Energy Appl ications at Low TemperaturesIV , NBSIR 81-1645, National Bureau of Standards, Boulder, Colorado

pp. 41 5-451 ( 1981 )

.

12. T. A. Whipple, H. I. McHenry, and D. T. Read, "Fracture Behavior of

Ferrite-free Stainless Steel Welds in Liquid Helium," in Material

s

Studies for Magnetic Fusion Energy Appl ications at Low Temperatures - III , NBSIR 80-1627, National Bureau of Standards, Boulder, Colorado,

pp. 167-194 (1980).

312
LIST OF TABLES

Table 1.

Table 2

Table 3.

Table 4.

Table 5.

All-weld metal composition and ferrite content.

Delta-ferrite measurements on double-vee 31 6L stainless steel

wel dments

.

CVN impact test results on 31 6L weldments.

Tensile properties of 316L stainless steel weldments.

Fracture toughness properties of 316L stainless steel weldments

at 4 K.

313
Table 1

All -Weld-Metal Composition and Ferrite

Element

316L Bare 316L Tubular

GTA GMA SA GMA SA

SA High-Heat

Input

C 0.017 0.017 0.023 0.014 0.016 0.017

Mn 1.39 1.40 1.97 2.17 2.50 2.43

Si 0.45 0.47 0.84 0.46 1.03 0.84

Cr 18.02 17.95 18.95 18.68 19.61 19.07

Ni 11.89 11.76 12.75 15.15 15.32 14.59

Mo 2.12 2.12 2.08 2.31 2.21 2.22

S 0.003 0.003 0.005 0.013 0.011 0.011

P 0.020 0.020 0.031 0.010 0.021 0.021

N 0.069 0.069 0.065 0.030 0.016 0.018

0 0.053 0.066 0.110 0.072 0.106 0.087

Co 0.092 0.090 0.095 0.012 0.025 0.020

V 0.063 0.063 0.06 0.013 0.017 0.016

Nb 0.02 0.02 0.02 0.02 0.02 0.02

Ti 0.021 0.022 0.024 0.15 0.092 0.104

Cu 0.089 0.086 0.10 0.037 0.043 0.042

FN measured 1.5 3.0 7.4 0.9 2.2 1.8

FN calculated

by diagram 4.4 4.1 5.7 0.8 6.8 5.2

314
Table 2

Del ta-ferrite Measurements on Double-vee 316L

Stainless Steel Weldments

Percent Del ta-ferrite*

Wire Process A B C D E Avg.

316L GTA 0.0 0.6 1.9 0.7 0.0 0.6

Bare

Wire GMA 3.5 3.0 2.8 0.6 0.0 2.0

316L GMA 0.5 3.2 2.3 3.2 0.6 2.0

Tubular

Wire SA

(high-heat

4.7 0.2 0.7 0.6 2.7 1.8

input)

*Measurements by ferrite meter at indicated locations.

315
Table 3

CVN Impact Test Resul ts^

’ 2 on 316L Weldments

316L Bare Wire 316L Tubular Wire

Temperature

,

SA High

K GTA GMA SA GMA SA Heat Input

172 102 114 49 58 24 45

(0.91) (1.17) (0.41) (0.91) (0.30) (0.53)

144 89 109 43 47 19 40

(0.99) (1.19) (0.41) (0.69) (0.18) (0.53)

116 75 92 37 39 18 33

(0.69) (0.99) (0.28) (0.43) (0.18) (0.36)

77 58 77 29 31 16 39

(0.46) (0.81) (0.05) (0.43) (0.08) (0.38)

^Results given are energy absorbed in joules and lateral expansion in (mm).

2A11 values are the average of three tests

315
Table 4

Tensile Properties of 316L Stainless Steel Weldments

Wire Process

Temperature,

K

Yield

Strength

,

MPa (ksi)

111 timate

Strength,

MPa (ksi)

Elongation,

%

Reduction

in Area,

%

316L GTA 295 430 602 32 55

Bare (62.4) (87.3)

Wire

4 807 1549 34 27

(117.0) (224.6)

GMA 295 355 554 45 55

(52.0) (80.3)

4 779 1510 37 27

(113.0) (219.0)

SA 295 441 587 36 54

(63.9) (85.1)

316L GMA 295 332 503 30 34

Tubular (48.1) (73.0)

Wire

4 608 1162 29 20

(88.2) (168.6)

SA 295 387 542 32 41

(56.1) (78.6)

SA 295 322 555 34 38

(high (46.7) (80.5)

heat

4 606 1122 21 18

input) (87.9) (162.7)

317
Table 5

Fracture Toughness Properties of 316L Stainless Steel Weldments at 4 K

J TIc,

r_ 3

O

i

—

kJ/m^ MPa /m

Wire Process (in*lb/in^) ( ksi /Tn)

316L Bare GTA 138

Wire 148

143 (818) 181 (165)

GMA 103

130

116 (666) 163 (148)

316L Tubular GMA 111 (635) 159 (145)

Wire

SA (High

heat input)

47 (270) 103 (94)

318
LIST OF FIGURES

Figure 1 Micrographs of locations C and E on the 31 6L bare-wire GTA

weldment, showing the variation in ferrite content through the

thickness of the weldment. KOH etch, 400X.

Figure 2 Micrographs showing the difference in the cleanliness of the

weldments (a) 316L bare-wire GTA weldment (b) 316L tubular-wire

SA weldment. Unetched, 200 X.

319
(b)

Figure 1. Micrographs of locations C and E on the 31 6L

bare wire GTA weldment, showing the variation

in ferrite content through the thickness of the

weldment. KOH etch, 400X.

320
(a)

(b)

Figure 2 Micrographs showing the difference in the

cleanliness of the weldments (a) 31 6L bare

wire GTA weldment (b) 316L tubular wire SA

weldment. Unetched, 200X.

321

A*

FRACTURE TOUGHNESS PROPERTIES OF WELDS IN ALUMINUM ALLOY 5083-0 AT 4 K

J. W. Elmer and H. I. McHenry

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

A series of weldments was prepared by five cooperating companies in

51-mm thick 5083-0 aluminum plate using 5183 filler metal and the gas

metal arc (GMA) welding processes. The J-integral fracture toughness

and fatigue crack growth rates of the weldments were measured at 4 K and

found to be similar to base metal properties. In a related study, the

tensile and notched tensile properties of the same weldments were obtained at room temperature and 4 K. Data for each of the mechanical

properties did not vary significantly from weld to weld. It is concluded that the mechanical properties of 5183 welds at 4 K do not depend

on the GMA welding parameters and procedures.

323

INTRODUCTION

Aluminum alloy 5083-0 is a solid-solution-strengthened alloy frequently used for cryogenic storage tanks for service to 76 K. Alloy

5083-0 derives its limited strength from the solid-solution-strengthening

contributions of 4.4% Mg and 0.7% Mn. It is readily weldable with alloy

5183 filler metal and has outstanding toughness at low temperatures

.

New applications for 5083-0 aluminum in magnetic fusion energy systems

are being examined; these place additional requirements on the alloy

since it is necessary to retain favorable mechanical properties down to

4 K. To date, few mechanical property data are available for 5083-0

weldments at 4 K.

In this study, fracture toughness and fatigue crack growth rate

data were obtained for 5083-0 weldments at 4 K. This study was preceded

by an evaluation of the tensile and notched tensile properties of the

same weldments at room temperature and 4 K by Kelsey and Mueller.^-

The objectives were as follows:

1. Develop fracture toughness and fatigue crack growth rate data

for 5083 aluminum weldments at 4 K.

2. Compare the 4-K toughness of 5183 aluminum weldments produced

by different companies using different welding parameters and

processes

.

WELDING PROCEDURES

Weldments were prepared and donated by five different companies.

Each company was asked to weld 51 -mm (2-in) thick 5083-0 aluminum plate

using 5183 filler wire and the gas metal arc (GMA) welding process. All

welding was performed in the flat position using a double-vee joint

325
preparation; however, specific welding parameters were the choice of the

individual company. In addition, two companies supplied high deposition

rate (HDR) GMA weldments of the same material to determine if high heat

input influences the mechanical properties. The welding parameters used

are listed in Table 1. The welding parameters varied considerably; for

example, the number of passes varied from 2 for an HDR weld to 55 for

one of the conventional welds. A total of seven welds were made, five

by the conventional GMA welding process and two by the HDR-GMA welding

process

.

TESTING PROCEDURE

All fracture toughness testing was performed in liquid helium at

4 K using a standard 2.54-cm (1-in) thick ASTM E399 compact tension

specimen. The specimen was removed from the center of the 51 -mm welded

plate with the plane of the notch perpendicular to the plate and centered in the fusion zone, such that the 5183 weld metal was tested. The

specimens were precracked in tension-tension fatigue to a depth of

3.05 cm (1.20 in) prior to testing. The fracture toughness was measured

using a computer-aided, single-specimen J-integral test, described by

f3l Read. J Fatigue crack growth rate measurements were made using the

same specimen configuration and were tested in liquid helium with an R

ratio of 0.1 .

RESULTS

Tensile properties of the 5083 weldments at room temperature

(Table 2) and at 4 K (Table 3) are provided to include additional information about the weldments.

^Table 3 also shows the 4-K fracture

toughness properties for each of the seven weldments as measured by the

326
J-integral technique. In Table 3, Kj

c (J) refers to the plane strain

fracture toughness as estimated from the J-integral test results.

Fatigue crack growth rate data at 4 K (Table 4) compares conventional

and HDR weld metal properties with 5083-0 base metal properties measured

F41

by Tobler and Reed. J The empirical constants n and C refer to the

Paris equation of fatigue crack growth rate:

C(AK)

n

C has metric units corresponding to da/dN in mm/cycle and aK in MPav/ m.

DISCUSSION

A comprehensive summary of the tensile data indicates that welding

parameters do not influence the yield strength or tensile strength to

any significant extent. Except for the generally lower tensile properties of the GMA weldment of supplier 2, the tensile properties of the

GMA weldments were quite consistent. The HDR weldments proved to be

only slightly lower in strength than the conventional weldments. The

notch yield ratios at 4 K and room temperature are all greater than 1.6,

which suggests that the weld metal is very tough.

Results of the fracture toughness testing at 4 K show the average

fracture toughness, Kj (J), determined by the J-integral method, to be

45.1 MPa/m. The K T (J) data are consistent for the seven weldments and

Ic

all the data lie within one standard deviation of the mean except for

that of the HDR weldment from supplier 4. The tensile data are also

consistent indicating that the mechanical properties of weldments in

5083-0 aluminum are relatively insensitive to welding parameters.

327
The average toughness of 45.1 MPa/m at 4 K is low when compared

with those of other tough materials used at low temperatures and seems

to be in conflict with the notch-yield-ratio results. Actually the

Kj

c (J) values for the welds are the same as the 5083-0 base plate (TL

[31 orientation) results at 4 K reported by Tobler and Read. J The low

values are due to the testing technique used and do not reflect the true

fracture resistance of the material. The problem with measuring the

toughness of the 5183 aluminum weldments using the J-integral technique

is that the aluminum tears at low J-levels where K T (J) is measured and

Ic

thus the K t (J) value is low. However, 5183 aluminum weldments are able

Ic

to sustain much higher J-levels under stable crack propagation conditions, which is an indication of toughness. This is shown in Figure 1,

which is a load-displacement curve recorded during one of the tests.

The point indicated is the point of initial tearing, determined from the

J-Aa curve. This occurs at a low load, but the specimen was able to

sustain loads past the peak load without fracturing. Thus, 5183 aluminum weldments are tough even though their Kj

c (j) values are low.

Fatigue crack growth rates for 5183 weld metal deposited by conventional and HDR-GMA processes are compared with the fatigue crack

growth rate of 5083-0 base metal at 4 K in Figure 2. The fatigue crack

growth rates for the conventional and HDR welds are similar and generally faster than the base metal rates; however, without further study

the differences between them cannot be considered significant.

SUMMARY AND CONCLUSION

Tensile properties at room temperature and at 4 K for 5183 aluminum

weld metal were shown to be consistent for welds made with large variations in welding parameters . ^ ^ The Kj

c (J) values and fatigue crack

328
growth rates at 4 K for the welds were also shown to be insensitive to

welding process and parameters. Therefore, it can be concluded that the

mechanical properties of 5183 weld metal do not depend upon the GMA

welding parameters to any great extent.

329
REFERENCES

1. R. A. Kelsey and L. N. Mueller, "Tensile and Notch-Tensile Properties at Room Temperature and 4 K of Welds in Aluminum Alloy 5083-0

Plate," in Material Studies for Magnetic Fusion Energy Applica -

tions at Low Temperatures-II

I

, NBSIR 80-1627, 219-244 (1980).*

2. F. G. Nelson, J. G. Kaufman, and E. T. Wanderer, "Tensile Properties and Notch Toughness of Groove Welds in Wrought and Cast Aluminum Alloys at Cryogenic Temperatures ," Advances in Cryogenic

Engineering, Vol . 14, 1969, pp. 71-82.

3. D. T. Read, "The Computer-Aided J-Integral Test Facility at NBS,"

in Materials Studies for Magnetic Fusion Energy Applications at

Low Temperatures-II

I

, NBSIR 80-1627, 205-216 (1980).

4. R. L. Tobler and R. P. Reed, "Fracture Mechanics Parameters for a

5083-0 Aluminum Alloy at Low Temperatures ," Journal of Engineering

Materials and Technology , Vol. 99, (October 1977).

*N0TE: A slightly revised version of Reference 1 (same authors and

title) has been published: Report Number 57-80-02, Alcoa

Laboratories, Alcoa Center, PA 15069, March 14, 1980.

330
Table 1

Wei d i ng Procedures [1]

Suppl ier

Process

Weld

Current

(A)

Vol tage

(V)

Ga s (L/min)

Number

of

Passes

1 GMA 230-250 26-29 15.5 Ar 55

31.0 He

2 GMA 250-300 26-29 11.7 Ar -- 35.2 He

3 GMA 240-300 24-26 28.2 Ar 25

4 GMA 335-365 33-34 9.4 Ar 16

18.8 He

5 GMA 380-390 28 — 18

4 HDR* 580 34 70.5 Ar

47.0 He 4

5 HDR 740-750 33 11.7 Ar

35.2 He 2

*High Deposition Rate

331
Table 2

Room-Temperature Mechanical Properties of 5083 Welds [1]

Suppl ier Process

0.2%

Yield

Strength

(MPa)

Notch

Tensi 1

e

Strength

(MPa)

Notch*

Yield

Ratio

1 GMA 166 312 1 .88

2 GMA 148 315 2.13

3 GMA 166 314 1 .89

4 GMA 167 327 1 .96

5 GMA 178 356 2.00

4 HDR 152 312 2.06

5 HDR 146 306 2.09

* NYR = Tensile Yield Strength

Notch Tensile Strength

Table 3

Mechanical Properties of 5083 Welds at 4 K [1]

Suppl ier Process

0.2 %

Yield

Strength

(MPa)

Notch

Tensi 1

e

Strength

(MPa)

Notch

Yield

Ratio

K,

C

( J )

(MPa/m)

1 GMA 202 363 1 .80 41 .7

2 GMA 177 344 1 .94 46.9

3 GMA 213 347 1 .63 41 .3

4 GMA 200 362 1 .81 42.8

5 GMA 214 391 1 .83 49.0

4 HDR 202 343 1 .69 51 .9

5 HDR 200 337 1 .69 42.0

332
Table 4

Fatigue Crack Growth Rate at 4 K

Material Condi tion n C AK

Range

MPav7^

5083-0 anneal ed 5.2 1 .3 ln -io

x 1 0 10-30

5183 HDR weld 6.5 3.8 in" 12

x 10 15-31

5183 GMA weld 4.6 7.3 x m-10 1 0 14-32

333
Initial Tearing

d>

m ‘avoi

0^1

J u~r

0

0.1

0.2

0.3

LOAD

LINE

DISPLACEMENT,

mm
TENSILE AXIS

50 ym

Figure 1 Micrograph of a 31 6L SA weld metal tensile

specimen deformed at room temperature, taken

near the fracture surface and showing the

elongation of the ferrite. KOH etch.

435

THE INFLUENCE OF FERRITE AND NITROGEN ON MECHANICAL PROPERTIES

OF CF8M CAST STAINLESS STEEL

Leroy Finch

ESCO Corporation

Portland, Oregon

ABSTRACT

The production and mechanical -property testing of a cast

molybdenum stainless steel are described.

337

THE INFLUENCE OF VARIABLE FERRITE AT A CONSTANT NITROGEN LEVEL

AND OF VARIABLE NITROGEN AT A CONSTANT FERRITE LEVEL ON

MECHANICAL PROPERTIES OF CF8M CAST STAINLESS STEEL

INTRODUCTION

The National Bureau of Standards in Boulder desired to study the effect of a) variable ferrite content at a constant nitrogen level and b) a variable nitrogen content at a constant ferrite level in cast molybdenum stainless steel, CF8M.

These materials were to be tested in tension at room temperature and

77°K (liquid nitrogen) at ESCO and at 4°R (liquid helium) at NBS in

Boulder. Impact resistance values were also to be determined at room temperature and 77°K at ESCO and at 4°K at NBS.

- —Tables 1 and 2 note the aim chemical analysis of the ten compositions

decided on, five with variable ferrite content from 0%, fully austentic,

to 28% with a nitrogen content of 0.05% and five with a calculated ferrite content of 10% with variable nitrogen contents from 0.02 to 0.20%.

The ferrite contents were calculated using the data and curve reported

at the recent ASTM symposium on cast stainless steels, "Ferrite Measurement and Control in Cast Duplex Stainless Steels" by L. S. Aubrey et al.

Figure 1 is the curve used to estimate the ferrite content from the

chemical compositions.

Pattern X3829 was used to make the molds for the casting to be used

by NBS for the tests to be made at 4°K. The final casting is 4"kl2"x24*.

The test bars tested at ESCO were from a conventional keel bar which has

a diameter of IV at the bottom.

PRODUCTION

Ten one thousand pound heats were melted in an induction furnace with

a monolithic magnesia lining and powered by a motor generator having a

960 Hertz output.

The sand mix used to form the molds was 33% zirconium silicate, 66%

silica bonded with bentonite. This was washed with an alcohol base

silica wash and torch dried. When the metal was all melted a preliminary

sample was sent to the laboratory for spectrographic analysis. A "suck-up"

sample was also sent for anlaysis of nitrogen content.

When the results of these analyses were received back at the melting

location, additions were calculated to bring the composition to the desired level. These additions were carefully weighed out and added to the

molten metal. When the proper temperature was reached as measured with

a platinum rhodium thermocouple, the metal was tapped into a magnesia

lined ladle and transported to the mold. Final chemical analysis samples

* were taken from the ladles full of metal. When the proper pouring temperature was reached the metal was ooured into the mold.

339
After the metal had cooled down the castings were removed from the

molds, heated to 2000 °f

,

water quenched and the risers burned off using

an oxyacetylane-powder torch. After the risers were removed the castings

were again heated to 2000°F and water quenched. This double heat treatment was used because it has been found that CF8M is susceptible to sigma

formation during cooling in the mold and if the risers are cut from material with sigma present in the microstructure cracking is likely to occur

in the areas subjected to thermal shock.

By heating to 2000°F the sigma is converted to ferrite and this ferrite

is retained by the rapid water quench. After the risers are removed the

material is again heat treated to redissolve any carbide which may have

precipitated during the removal of the riser and also to dissolve any

sigma that may have formed.

Conventional tension test bars were machined from keel bars poured with

metal from each heat. These were heat treated with the castings. Charpy

V notch impact test specimens were also machined from the keel bars. The

ferrite content of each heat was calculated using the Schoefer ratio and

curve, Figure 1. It was also measured with a calibrated Magna gage and

a Ferr itescope . A portion of the end of each of the tension test bars

was polished, etched and a point count made of the ferrite present using

the procedure given in ASTM-E 562-76.

The spectrographic chemical analyses of the heats produced are listed

in Tables 3 and 4. The nitrogen contents included in the tables were

obtained by vacuum fusion using a Leco rapid nitrogen analyzer. Also

included in the tables are the ferrite contents determined by the methods

noted above. The heat which was designed to have 0.02% nitrogen was melted

in an argon atmosphere in order that the metal would not absorb nitrogen

from the air and finish with a nitrogen content higher than 0.02%. The

ferrite contents of four of the original heats were lower than desired.

They were not used. The nitrogen content of another was higher than desired so it also was not used.

Five more heats were produced in a similar manner as described above.

In order to insure that the ferrite contents of these were satisfactory

before the molten metal was tapped the spectrographic sample poured into

the copper chill mold was measured with the calibrated Ferritescope

.

When the ferrite content was that desired the metal was tapped and poured

as before.

340
MECHANICAL PROPERTIES

Table 5 notes the room temperature mechanical properties obtained at

ESCO for the ten heats which will be tested by the NBS . Table 6 lists

the impact properties obtained to date. The tensile properties of the

heats have not been obtained at 77°K at this time nor have the impact

tests been completed for the last four heats. They will be reported

when obtained. Figure 2 is a plot of the tensile and yield strengths

(at 0.2% offset) at room temperature of the variable ferrite content

heats versus the ferrite content. The ferrite content is expressed as the Cr^/Ni 1 ratio. Figure 3 is the plot of the tensile and yield strengths

(at 0.2% offset) at room temperature of the variable nitrogen content heats versus the nitrogen contents.

In Figure 3 the numbers above each point plotted show the ferrite content of the material expressed by its ratio. The variable nitrogen heats

did not have exactly the same ferrite contents. Heat 48987 was not used

in these plots since the ferrite was 14% while the others were 8-10%.

Heat 48988 with 8% ferrite and 0.05% nitrogen was used.

Figure 4 is a plot of the average impact strength at room temperature

and 77 °k versus the ferrite content. Figure 4 is a similar plot of

impact strength versus nitrogen contents.

DISCUSSION

From Figure 2 and Table 5 it can be said that the ultimate tensile

strength increases 940 psi per 1% ferrite when the ferrite varies from

0%, fully austentic, to 29%. The yield strength increases 890 psi per

1% ferrite over the same range of ferrite.

From Figure 3 and Table 5 it can be said that increasing the nitrogen

content from 0.02 to 0.05% has little or no effect on mechanical properties. When the nitrogen content increases from 0.05 to 0.14% the tensile

strength increases 855 psi per 0.01% of nitrogen. When it increases from

0.14 to 0.20% the increase in tensile strength averages 1250 psi per 0.01%

of nitrogen.

From Figure 2 and Table 6 it will be noted that the increase in tensile

•strength between 8 and 10% ferrite, 1.14 and 1.19 ratio is 2100 psi so if

2100 is subtracted from 93.7 the tensile strength of the heat with 1.19

ratio and 0.20% nitrogen the tensile strength of the material if it had

had 8% ferrite and 0.20% nitrogen could have been 91.6 ksi, so 91.6 - 84.1 = 7500 psi, which could have been due to the increase in nitrogen

content or 1250 psi per point of nitrogen.

341
——It will be noted with similar calculation that the tensile strength

of the heat with 0.10% nitrogen would have fallen on the line if its ferrite content had been 8%, 1.14 ratio, instead of the 10%, 1.17 ratio.

The effect of the nitrogen content increase on yield strength appears to

be linear, „ 1160 psi per 0101% nitrogen. It will be noted in Figure 4 that the impact strength decreases as ferrite content increases. This

is true at room temperature as well as at 77°K. Figure 5 shows that the

impact strength at room temperature and 77°K is also lowered by increasing

the nitrogen content. Of course this would be expected with the tensile

strengths increased.

When the mechanical property data is completed the results will be

reported.

342
Table 1 - Aim Chemical Analysis in Percent of Five Heats of

CF

8

m with Variable Ferrite and Constant Carbon and

Nitrogen Contents

Ferrite % 0 7 14 21 28

C 0.06 0.06 0.06 0.06 0.06

Mn 1.20 0.40 0.40 0.40 0.40

Si 0. 50 1.25 1. 25 1.25 1. 25

Cr 18.25 18. 35 18.65 19.20 20. 50

Ni 13.00 10.75 9.05 8.20 8.20

Mo 2. 15 2. 15 2.15 2.15 2.15

N

2 0.05 0.05 0.05 0.05 0.05

Cr 1 /Ni 1 0.90 1.12 1.27 1. 39 1.485

Table 2 - Aim Chemical Analysis in Percent of Five Heats of

CF8M with Variable Nitrogen Contents at a Constant

Ferrite Level, 10%

C 0.06 0.06 0.06 0.06 0.06

Mn 0. 40 0.40 0.40 0.40 0.40

Si 0.75 1.25 1. 25 1. 25 1. 50

Cr 18. 55 18. 90 19.25 21.20 21.00

Ni 10.10 10.00 9.25 9.45 9.25

Mo 2.15 2.15 2.15 2.15 2. 85

N

2 0.02 0.06 0.10 0.15 0.20

Cre '/Nie ’ 1. 19 1.19 1. 19 1. 19 1.19

343
Table 3 - Final Spectrographic Analysis in Percent of Five Heats of

CF8M with Variable Ferrite Content and a Constant Carbon

and Nitrogen

Heat No. C Mn Si Cr Ni Mo W

2 Ferrite Content*

47640 0.06 1. 10 0.59 18.04 13.24 2. 12 .05

C7¥8

M

0 Pt Ct0 48988 0.06 0.38 1. 16 19.62 10.61 2,12 .05 1. 21 6 8 47521 0.06 0. 36 1. 13 19.23 8.25 2.16 .05 1. 39 17 15

47552 0.06 0.35 0.95 20.40 8. 22 2.13 . 05 1.45 22 22

49021 0.05 0. 39 1.11 22.54 9 .84 2.26 . 05 1. 48 27 29

Table 4 - Final Spectrographic Analysis in Percent of Five Heats with

Variable Nitrogen at a Constant Ferrite Content

Heat No. C Mn Si Cr Ni Mo N

2 Ferrite Content*

49020 0.07 0.39 1. 09 19.32 10.00 2.07 .02

C M

1775 8 Pt Ct

8 48987 0.07 0.51 1. 17 20.13 10.23 2.49 .06 1.27 11 14

47364 0.06 0.40 1. 30 19.17 9.14 2. 00 .10 1. 18 8 9 47390 0.07 0.41 0.98 21.00 9.46 2.13 . 144 1. 16 7 8 47399 0. 06 0. 33 1. 35 21.66 9.68 2. 83 . 20 1. 19 9 10

*Ferrite Content

C - Calculated using Schoefer formula for ratio

Cr

1 =Cr+l. 5Si+l. 4MO-4 . 99

Ni 1 = 30xC+ . 5Mn+Ni+26 (N-02) +2. 77

M - Measured with calibrated Ferritescope

Pt Ct - Metallographic point count, ASTM E 562

The Cre'/Nie' ratios used in the plotting of the figures were obtained

from the ferrite content determined by point count and taken from

Figure 1.

344
Table 5 - Room Temperature Mechanical Properties of Heats Shipped

to NBS Tested at ESCO

Heat No.

UTS

ksi YS ksi E (0.20% offset ksi ) % RA

% Ferrite

Variable Ferri te , 0 . 0 5% N Ratio %

47640 65. 3 29 .

6

55. 5 65 .9 0

48988 76. 4 36.4 56.0 71 1. 14 8

47521 87.2 48.5 48. 0 64 1. 30 15

47552 89 .

0

50.2 46. 5 63 1. 43 22

49021 92.6 55.4 42.0 64 1. 53 29

Variable Nitrogen Content,. Constant Ferrite

49020 76.

4

36.2 58. 0 65 1. 14 8 48987 82.9 41.0 47.0 69 1.28 14

47364 81. 3 44. 3 48.5 70 1. 17 9 47390 84. 1 47.6 51.0 66 1. 14 8 47399 93.7 53. 8 42. 0 59 1.19 10

Table 6 - Impact Properties Obtained to date

Impact Data,, Ft-Lbs

Heat No. R. T. Avg 7 7°K

.

Avg

.

47640 177.-240-240 219 136-122-147 135

489 88

47521 156-140-151 149 74- 78- 76 76

47552 149-122-145 139 63- 66- 61 63

49021

49020

48987

47364 167-232-223 207 66-101-110 92

47390 170-175-169 171 65- 68- 61 65

47399 190-175-172 179 38- 39- 22 33

345
ADDENDUM

The Influence of Variable Ferrite at a Constant Nitrogen Level

and of Variable Nitrogen at a Constant Ferrite Level

on Mechanical Properties of CF8M Cast Stainless Steel

This is an addendum to the original report which did not contain

all of the data.

This report includes the missing data: impact test results at room temperature and 77°K and tensile properties at 77°K.

Table 6 reports the impact strength at room temperature and 77°K

obtained on all of the alloys. Some of these were included in the

original report.

Figure 4 is a plot of the impact strengths of the heats with

variable ferrite content vs. the ferrite content at room temperature

and 77°K.

Figure 5 is a plot of impact strength vs. the nitrogen contents

o for the variable nitrogen series at room temperature and 77 K. The impact strength drops with increase in ferrite content at

both room temperature and at 77°K.

At room temperature the loss per percent of ferrite is as follows

0 - 8 %

3.9

ft- lbs

0-15% 4.7 ft- lbs

0 - 22 % 3.6 ft- lbs

0-29% 3,4 ft- lbs

At 77°K the drop is:

0 - 8 %

3.9

ft- lbs

0-15%

3.9

ft- lbs

0 - 22 % 3.3 ft- lbs

0-29% 3.1 ft- lbs

The rate of loss is nearly equal at both temperatures.

Table 7 reports the tensile properties obtained at 77°K.

346
Figure 6 is the plot of the tensile strengths at 77°K plotted

against the ferrite content expressed as Cr^/Ni^ ratio and the

nitrogen contents.

It will be noted that the tensile strength increases both with

ferrite content and with nitrogen contents to maxima and drops off

with further increases: 22% ferrite with a ratio of 1.43 has the

greatest strength of the ferrite series and the nitrogen content

of 0.10 percent was the strongest of the nitrogen series of heats.

LFinch

/nmo

2/25/81

347
Table 5 - Room Temperature Mechanical Properties of Heats Shipped

to NBS Tested at ESCO

Heat No.

UTS

ksi (

0

YS ksi

.20% offset

E ksi ) % RA

% Ferrite

Variable Ferrite, 0.05% N Ratio % 47640 65.3 29.6 55. 5 65 .9 0 48988 76.4 36.4 56.0 71 1. 14 8 47521 87.2 48.5 48.0 64 1. 30 15

47552 89 .

0

50.2 46. 5 63 1.43 22

49021 92.6 55.4 42.0 64 1.53 29

Variable Nitrogen Content, Constant Ferrite

49020 76.4 36.2 58. 0 65 1.14 8 48987 82.9 41. 0 47.0 69 1.28 14

47364 81. 3 44. 3 48.5 70 1. 17 9

47390 84. 1 47.6 51.0 66 1.14 8 47399 93.7 53. 8 42.0 59 1.19 10

Table 6 - Impact Properties

Impact Data, Ft-Lbs n Heat No. R.T. Avg

.

77°K Avg

.

47640 177-240-240 219 136-122- 147 135

48988 203-162-199 188 98-108- 107 104.

47521 156-140-151 149 74- 78- 76 76

47552 149-122-145 139 63- 66- 61 63

49021 125-118-115 119.3 40- 45- 47 44

49020 162-137-139 146 73- 87- 85 81.

'

48987 184-198-205 195.7 118-102- 93 104.:

47364 167-232-223 207 66-101- 110 92

47390 170-175-169 171 65- 68- 61 65

47399 190-175-172 179 38- 39- 22 33

348
Table 7 - Mechanical Properties at 77°K of Heats Shipped to NBS

Tested at ESCO

Ferrite

Heat No. UTS, psi E, % RA, % Ratio _%

Variable Ferrite Content, 0.05% N

47640 142,300 38 47 . 9 0 48988 190,800 53 54 1.14 8 47521 211,200 30 20 1. 30 15

47552 214,100 37 28

• 1.43 22

49021 169,200 30 56 1. 53 29

Variable Nitrogen Content, Constant Ferrite N

2 49020 183,100 35 55 1. 14 8 .02

48988 190,800 53 54 1.14 8 .05

48987 186,000 53 61 1.28 14 .06

47364 202,700 35 25 1. 17 8 .10

47390 173,300 39 52 1. 14 8 . 144

47399 170,500 25 36 1. 19 10 .20

LFinch

/nmo

1/23/81

349
Cr'

/Ni

Cr

1 = %C R + 1 . 5%S i = 1 .4%Mo - 4.99

Ni

1 = 3 0 ( %C ) + + %Ni + 26(%N - .02) + 2.77

Figure 1 Relationship between Cr^/Ni^ and percent Ferrite

in Cast Duplex Corrosion Resistant Stainless Steels

350
%CR

+1

.5%Si

+1

.mo

-

4.99

is>i ‘HiONaais aiaiA

r-^ CM

CMO

CO

CM

CM

O00

I! II

S- 'I— o ^

|S>| ‘H10N3dlS 3HSN31

351

Figure

2

Relationship

between

RT

Tensile

and

Yield

Strength

and

Ferrite

Content
!S>J ‘H10N3H1S 0131

A

o o o o

!S>) ‘H10N3U1S 31ISN31

+-> cn

O)

s_

4-> <T> -a

'a>

<u

cn

35

NITROGEN

CONTENT,
3

(0

%C)

+^

+

%Ni

+

26

(%N

-

.02)

+

2.77

X ZZ ‘sqi-U ‘H19N3U1S JLOVdIAII

~a

S- •(- o zr

la ‘sqi-W ‘HlONiaiS lOVdlAII

353

Figure

4

Relationship

between

Impact

Strength

at

aii

77K

and

Ferrite

Content

expressed

as

Cr

/Ni
220

>1 LI ‘sqi-W ‘H1DN3U1S iOVdlfll

1U ‘sqi-W ‘H10N3U1S IDVdMI

NITROGEN

CONTENT,

%
FERRITE

CONTENT,

Cr'/Ni

0.8

1.0

1.2

1.4

IS)| ‘H10N3U1S 31ISN31 31Vlflll~in

355

NITROGEN

CONTENT,

%

Figure

6

Relationship

between

Tensile

Strength

at

77

and

Ferrite

content

expressed

as

Cr’/Ni

1

.

Cr

=

%CR

+1

.5%Si

+1

.

4%

Mo

-

4.99

Ni 1

=

30(%C)

+

~

+

%Ni

+

2S(

%N

-

.02)

+

2.77

METALLOGRAPHY OF 5 -FERRITE IN AUSTENITIC

STAINLESS STEEL CASTINGS

Elliot L. Brown

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

The volume fraction and morphology of 6-ferrite in austenitic

stainless steel castings and welds profoundly influence their mechanical

behavior over a wide range of temperatures. A review of the literature

summarizes the effects of composition, solidification, and solid state

(5->y) transformation on the resultant 6-ferrite microstructure in welds

and castings. An optical metallographic study was performed on three

ASl'M A351 Grade CF8M alloy castings (e.g. 18-21Cr, 9-1 2Ni , 2-3Mo)that

possessed different 6-ferrite contents. Metallographic techniques

permitted distinction between the solidification and 6->y transformation

behavior in these castings. All of the alloys solidify in a primary 6 mode in which y precipitates before solidification is complete. The

residual 6-ferrite microstructure, owing to the 6-»y transformation, is

generally characterized by curved segments of 6/y interface connected by

straight or faceted segments. Frequently occurring interfacet angles

are consistent with the facets being parallel to a { 1 1 0 } 6 crystallographic plane and an orientation relationship between 6 and y across the

faceted 6/y interface. The metallographic features of the 6-ferrite

microstructure can be rationalized with the aid of a 6+y transformation

model analogous to that proposed for the proeutectoid y-m transformation.

357

INTRODUCTION

The solidification behavior of austenitic stainless steels is of

concern in both welding and casting technology. The term "austenitic"

stainless steel is actually a misnomer with respect to weldment and

casting microstructures , since they may contain significant amounts of

del ta-ferrite (6-ferrite), the high temperature bcc phase of iron, as

well as the fee austenite phase (y). The presence of 6-ferrite can

significantly alter the physical and mechanical properties of austenite.

Property variations due to 6-ferrite can be beneficial or deleterious

depending upon the nature of the material application and its temperature

range. For example, in the case of weldments the presence of 5 to 10%

1 2 6-ferrite has been reported to reduce the incidence of hot cracking ’

and microfi ssuri ng . At high temperatures (>973 K) 6-ferrite can transform

to the brittle sigma (a) phase, which results in a degradation of mechanical

d- f) properties . In highly corrosive environments a-phase can provide a

path for selective corrosion attack

7 of welds and castings. Delta-ferrite

in the austenite matrix al so • generally increases the strength and lowers

O fracture toughness of welds and castings , especially in cryogenic

g applications .

Since 6-ferrite profoundly affects the properties of austenitic

stainless steel, it would obviously be advantageous to control its

volume fraction and morphology in welds and castings. This report will

consist of a review of 6-ferrite in austenitic stainless steel followed

by a description and discussion of metal lographic studies of 6-ferrite

in stainless steel castings.
REVIEW OF 6-FERRITE IN AUSTENITIC STAINLESS STEEL

Composition Effects

Most research has been conducted in the area of welding austenitic

stainless steels and, more specifically, on the effects of various

filler metal compositions on 6-ferrite content. Some additional work

has been performed to extend the concepts developed for weld metals to

castings. Quantification of the effects of composition on 6-ferrite

content are conveniently accomplished by Cr- and Ni-equi val ents . The

Cr-equi val ents integrate the effects of ferrite formers (Cr, Mo, Si) and

the Ni-equival ents , the effects of austenite formers (Ni, Mn, C, N).

Delong^ reviewed ferrite in austenitic stainless steel weld metal and

traced the development of the Schaeffler and DeLong Diagrams (Figure 1).

Both of these diagrams relate 6-ferrite content, in volume percent

(v/o), to Cr- and Ni-equi val ents . The Schaeffler Diagram was determined

by quantitative metal lographi c techniques to within a ±4 v/o accuracy.

The 6-ferrite contents of the DeLong Diagram were determined magnetically

to within ±3 v/o, and a factor to account for the strong austenite

forming effect of nitrogen was added to the Ni-equival ent. The DeLong

Diagram predicts higher 6-ferrite contents than the Schaeffler Diagram

for more highly alloyed materials. A recent modification of these

diagrams has been made to account for the effects of precipitation in

stabilized austenitic stainless steels, since precipitation removes

carbon or nitrogen or both from solid solution, thus effectively reducing

1 3 austenite stability. Despite the fact that cooling rates encountered

during common welding practice do not exert a significant effect on

total ferrite content, DeLong cautions against a literal translation of

360
data on welds to castings."* Similar empirical correlations between 0

14 15

6-ferrite content and steel composition have been established for castings ’

and are contained in the Schoefer Diagram ( Figure 2). This diagram is

actually derived from the Schaeffler Diagram and is supported by measurements

made on castings."*

4

All of these diagrams refer to room temperature 6-ferrite formed in

a nonequilibrium manner. In fact, the equilibrium microstructures of

these austenitic stainless steels at room temperature do not contain any

6-ferrite. The departure from equilibrium can be effected by altering

the solidification rate, the cooling rate subsequent to solidification,

or the chemical composi tion

."* 0 The following sections discuss the

specific ways in which this departure from equilibrium affects the

extent and morphology of 6-ferrite in austenitic stainless steels.

Solidification Effects

Numerous investigations have been conducted to elucidate the sol idifi17-26 1627

cation behavior in austenitic stainless steel welds and castings. ’

Basically, the solidification behavior of this class of stainless steels

can be divided into primary 6 and primary y modes according to the ratio

of Cr- and Ni-equivalents. This division is frequently illustrated by

construction of appropriate constant Fe pseudobinaries (isoplethes) of

17 18 27

the Fe-Cr-Ni ternary. 5 5 It is important to realize that pseudobinary sections through ternaries are only qualitative, inasmuch as tie

lines drawn across multiphase fields do not represent true equilibrium

3G1
28 tie-lines. They are useful, however, for qualitatively illustrating

solidification modes as a function of composition.

A typical pseudobinary of this sort is depicted in Figure 3. The

two-phase, y+L and 6+L regions are separated by a eutectic triangle of

y+6+L region. Below the triangle, a y + 6 phase field exists, enclosed

by the y and 6 solvus lines. The y

+ 6 region broadens with decreasing

iron content. Chromium- and Ni-equival ents are substituted for Cr and

Ni to increase the generality of the ensuing review. If the nominal

composition of the steel possesses a low Cr-equival ent (low proportion

of ferrite formers) the primary solid phase is austenite. Similarly, in

a steel with a high Cr-equival ent (high proportion of ferrite formers),

ferrite is the primary solid phase. Therefore, solidification can be

divided generally into primary y and primary 6 modes. The solidification

mode becomes more complex in the vicinity of the eutectic triangle. For

nominal compositions C^ > Cr^ > C-j , the primary solid phase is still

austenite; however, ferrite precipitates before solidification is complete.

Likewise, for nominal compositions C^ > Ni > C^, the primary solid

phase is ferrite but austenite precipitates prior to complete solidification

The microstructures that result from solidification and cooling in the

vicinity of the eutectic triangle are also more complex. The classification

of solidification modes has been variously described in the 1 iterature.^*^

Table I summarizes these classification schemes. Composition ranges and

application (welding or casting) are included.

The ternary eutectic behavior described above is applicable to the

majority of commercial austenitic stainless steel grades. In the materials

22,27,29

352
under consideration the eutectic has a divorced morphology since the

23

amount of liquid at the eutectic composition is generally small. In

a divorced eutectic mixture one of the solidifying phases is discontinuous

and present in smaller volume fraction (i.e. embedded in the continuous

3(1 31 phase). * In primary y alloys the divorced phase at the subgrain

and grain boundaries is 6. In primary 6 solidified alloys the divorced

23

phase is y.

Studies of solidification in conventional ingots and by a directional

solidification and quenching technique have identified three possible

16 18 27 29 32

stages of solidification in an initially primary 6 alloy. * * *

*

.

Initial precipitation of 6 is interrupted by precipitation of y, and at

the end of solidification, a transition back to precipitation of 6 is

32

observed. This corresponds to the Type B solidification mode depicted

in Table I, L-*L+S-d_+ 6 +y->6 +y . In a conventional casting, the L+6+Lt5+y

transition has been associated with the transition from a columnar or

branched columnar solidification structure to an equiaxed structure.

Segregation of alloying elements ahead of the solidification front as

well as differences in cooling rate have been proposed as reasons for

the interruption of initial solidification of 6 by precipitation of y

from the melt and the transition back to 5 precipitation at the end of

1 71

solidification. The classification scheme of McTighe and Beech ’

(see: Table I) will be used in the following descriptions. Schematic

diagrams of Type A and Type B solidification extracted from the work of

these investigators are shown in Figure 4. For Type A alloys, the

primary 6 dendrites grow out into the liquid and 6 occupies close to 100

v/o of the as-solidified structre. With falling temperature y forms by

the solid state S-vy transformation, first in the primary grain boundaries

363
(y-j) and then intragranularly at larger undercoolings (y 2 and Y

3

)

• The

i ntragranul ar y can possess more than one crystal lographic orientation

(e.g., difference between y

2 and Y3 K For Type B alloys, the primary

6 dendrites grow into the liquid and a y envelope precipitates directly

onto the 6 from the liquid (y-j). The S-*y transformation results in the

growth of y (y 2

) from the envelope into the 6 dendrite. Some 5 remains

in the dendrite cores. The orientation of the 6 is the same within a

given grain or dendrite. Similar behavior has been observed by other

29 investigators.

The col umnar-to-equiaxed grain morphology transition often observed

in conventional castings takes place in Types A and Type B solidification.

However, in Type A alloys it is observable in both the primary and

secondary structures, whereas in Type B alloys it is only observable in

27

the primary structure (Figure 8 ). Primary structure refers to the

structure produced by solidification; secondary structure refers to that

33 produced by the subsequent 6 +y transformation. The primary and

secondary structures in Type B alloys can be differentiated by metal 1 ographic ... . , . 16,27,32,34 etching techniques.

Quenching techniques provide a powerful tool for analyzing the

course of solidification by allowing the isolation of a particular stage

of sol idification.^ s ^»27,29,32 By quantitative metal lography of

specimens quenched from various points on a cooling curve, the volume

fraction of primary 6 has been determined as a function of time or

29

temperature and composition for a cooling rate of 20 K/min. At this

cooling rate the maximum volume fraction of primary 6 was 0.5 prior to

the start of austenite growth in Type B alloys, and decreased monotonically

with increasing C- equivalent (C + 0.65 N) until the onset of primary y

freezing. The maximum volume fraction of primary 6 could be increased

354
further by alloying in the Type A alloy region (high Ni -equi val ent)

.

Other investigators have reported the ability to retain increasingly

larger amounts of residual 6 (subsequent to 6->7 transformation) by

16 21 32 35 increasing the cooling rate. * ’ ’ The observed increase in residual

6-ferrite may reflect, at least partially, an increase in primary 6 with

increasing cooling rate. This has been attributed to an increase in

16 27

departure from equilibrium conditions with increasing cooling rate. ’

In contrast, work done on austenitic stainless steel arc welds has

demonstrated only a small effect of cooling rate on total ferrite content,^

1 0 24

but a significant effect on ferrite scale and morphology. ’

Concomitant with the evolution of a solidification structure is the

redistribution of solute elements, commonly referred to as solute segregation.

Reference to the equilibrium phase diagram indicates the various possibilities,

36“ 38 which have been described in detail elsewhere. ” Under conditions of

equilibrium (slow cooling), diffusion in the solid and liquid and complete

mixing in the liquid allow for complete solute redistribution, so that

the composition of the solid formed is always that defined by the tie-line

intersection with the solidus curve. The solid composition at the

solidus temperature is the nominal composition. A more realistic situation

is one in which solid-state diffusion is minimal so that coring or

37 38 microsegregation occurs during solidification. ’ Perfect mixing in

the liquid and mixing in the liquid by only diffusion represent extreme

situations. Real cases probably exist between these two extremes. For

the cases described above, the final distribution of solute after normal

freezing is depicted in Figure 5. The theoretical analyses of segregation

are easily applied to planar interfaces, but it has already been implied
that more complex growth forms, such as cellular, cel 1 ular-dendritic,

and dendritic, are the case in most castings and welds. These solidliquid interface types are mainly due to the constitutional supercooling

3 6

produced by solute redistribution. However, they do not affect the

generality of the statements made with respect to planar interfaces

since the cell or dendrite can be viewed as composed of a series of

small volume elements that grow into the liquid as a planar interface,

A generalized theory of solidification of austenitic stainless

17

steel as primary 5-ferrite has been formulated for weldments by assuming

mixing in the liquid by diffusion only.^ A schematic diagram in

Figure 6, extracted from this work, depicts the solute distributions of

Ni and Cr across a subgrain following solidification. The gradients in

composition are associated with the subgrain (dendrite) core and the

subgrain boundary ( interdendritic) region.

Attempts to analyze the solute distribution in these alloys have

relied on electron microprobe and scanning transmission electron microscopy (STEM) analyses. The STEM possesses a higher spatial resolution

capability. Analyses have been performed on actual welds^’^’^

- ^ 5 ^’^

and specimens in which solidification has been interrupted by quenching

In conventional welds and castings difficulties arise in separating the

parts of the total solute redistribution due to solidification and the

6-vy transformation. Assumptions regarding the nature of solidification

and the 5->y transformation are generally necessary; e.g., no solute

17 24 25 redistribution during the 6-*y transformation. ’

’ Reasonably independent assessments of the solute distribution coefficients and the solute

profiles due to soldification have only been made in interrupted solidification studies.

18,29,32

3G6
Solid State 6->y Transformation and Residual 6-ferrite Morphologies

As-cast and as-welded austenitic stainless steel microstructures do

not generally represent an equilibrium state. Treatments that promote

equilibrium tend to dissolve 6-ferrite, e.g. wrought grades possess

virtually no 6-ferrite at room temperature and the 6-ferrite levels in

castings and welds can be substantially reduced by annealing at temperaO Q

tures of approximately 1400K. 9 A primary 6 alloy may contain in

excess of 50 v/o 6 upon solidification. Since the room temperature

microstructure of such an alloy contains substantially less 6-ferrite,

it is obvious that the 6+y transformation plays a major role in the

evolution of microstructure in duplex austenitic stainless steel welds

and castings.

A 6->y transformation only occurs for alloys that undergo primary 6

solidification, and the nature of the transformation appears to vary

with composition, cooling rate, and perhaps with solidification behavior.

It is important to note that a number of labels have been associated

with the various observed residual ferrite morphologies that are descriptive

in nature but may not describe the source of the morphology accurately.

This section attempts to describe the various ferrite morphologies, and

associate them with labels commonly employed in the literature.

Primary 6 alloys can be classified further into alloys remote from

the eutectic triangle that solidify with almost 100 v/o 6 (forming y

i nterdendri tical ly) and alloys that intersect the eutectic triangle, in

which solidification to y occurs directly from the melt after some 6 has

solidified. In the former type of alloy, the resultant room temperature

microstructure is completely determined by the solid state transformation.

The character of the phase transformation depends on the degree of

1 fi 17 undercooling, which in turn, depends upon the cooling rate. ’ For

357
relatively rapid cooling, the 6->y transformation has been referred to as 18 23 24 3? lq ??

either a Widmanstatten ’ ’

’ or an acicular * transformation,

but the former is probably a more accurate description. The Widmanstatten

phase transformation can be described as a thermally activated and

diffusion controlled reaction in which an orientation relationship

40

exists between the parent and product phases. In this case, the

19 20 22

residual 6-ferrite has been termed 'lathy

1

.

5

’ Although the residual

6-ferrite morphology has been termed 'lathy 1 , it is actually the austenite

that forms laths by the 6->y transformation. Microstructural investigation

20 of the lathy morphology in welds has indicated the presence of several

differently oriented lath bundles in a given large growth unit. The

model for Type A alloys, depicted in Figure 4a, describes this case.

The association of lathy 6-ferrite morphology with a Widmanstatten

transformation on quenching allows for the metal lographic delineation of

1 q 22 regions in a specimen that have solidified as primary 6. ’ At slower

1 9 22 cooling rates the mode of 6^y transformation has been termed 'equiaxial'. ’

It is not exactly clear what 6-ferrite morphology this type of 6->y

transformation produces.

Many primary 6 alloys of interest solidify so that y envelops the

primary 6 prior to completion of solidification. This mode is depicted

in the model of Figure 4b and refers to the sequence L-^L+6->L+6+y

.

In this case there is already solid y present in the microstructure in

the form of the y envelope, as a result of a L+y transformation. The

solid y envelope allows for heterogeneous nucleation of y during the 6-*y

transformation with less undercooling (higher temperature) than for the

Widmanstatten 6->y transformation. Therefore, the residual 6 exhibits

368
The morphology r 4-u w • j . , , -i 1 6,18,1 9,22,27,32

less of the Widmanstatten type morphology.

of the residual 6-ferrite in this case has been termed 'vermicular'

17 22 23

because of its roundish appearance.

5 ’ The residual 6-ferrite

morphology in primary 6 alloys has also been described as a cel 1-1 ike

22-24 27

network. ’ It is unclear whether this morphology is associated

with the vermicular or lathy 6 morphologies or is a transition morphology

between the two. A continuous spectrum of ferrite morphology, from

22

vermicular to lathy, has been observed with increasing Cr^/Ni^.

Regardless of the observed morphology, the residual 6-ferrite in

primary 6 alloys is definitely located at the cores or spines of dendrite

arms (subgrains). The progressive growth of y into the

primary 6 (see Figure 4) results in segregation of ' ferritizers ' to the

6 at the dendrite or subgrain core, thus stabilizing it against trans18 23 24 32 formation. ’ ’ * The formation of a y envelope occurs with attendant

segregation of alloying elements. Since diffusion, and hence homogenization,

of alloying elements is more difficult in austenite the segregation

persists. The segregation of alloying elements permits etching and

27

hence, observation of the prior dendritic form at room temperature.

In alloys that solidify as virtually 100 v/o 6, the ease of diffusion in

the bcc lattice minimizes segregation at high temperature, and therefore

the prior dendritic structure is not observable. Transmission electron

microscopy ( TEM) observations have shown that lathy 6 possesses a

Kurdjumov-Sachs (K-S) type orientation relationship^ with austenite,

whereas vermicular 6 possesses no crystal lographic relationship with

22 austenite. Angular faceting of 6/y interfaces in primary 6 alloys has

23 also been observed and noted in castings.

369
The observations of primary 6 alloys described above have all

18 22 implied a diffusion-controlled mechanism for the 6->y transformation. ’

Other investigations, in welds, suggest that the rapid cooling rates

attendant with relatively low heat input welding precludes the high

temperature residence time necessary for a diffusion controlled transformation. * ~ They propose, instead, a massive phase transformation

that would allow the bcc->fcc crystal structure transition to occur at

constant composition. The massive transformation is a thermally activated,

42 1

7

partitionless phase transformation. Lippold and Savage have proposed

that most of the 6 dendrite solidifies at constant composition (Figure

6). These workers are suggesting that this region of constant composition

17 23 29

6 transforms to y by a massive transformation. 5 5 The assertion

that the Ni and Cr segregation profiles characteristic of solidification

are preserved when a massive 6->y transformation occurs is supported by

quantitative chemical microanalysis (STEM) across a dendrite. The

concentration profiles measured are also cited as evidence for short-range

diffusion across the interphase boundary as the mechanism of transformation

(i.e., local depletion of Cr in y and excess in 6). Other investigations

have shown that a significant amount of microsegregation results from

18 29 32

the solid state 6->y transformation. ’

* Thus, it is uncertain

whether all the microstructural observations of residual 6-ferrite in

these investigations are consistent with or peculiar to the occurrence

of a massive transformation. Perhaps a massive 6->y transformation

occurs only under conditions of rapid cooling. It is possible that an

alternative explanation of the observed concentration profiles in welds

lies in the enhanced homogenization associated with finer dendrite

spacingsJ 6,27

370
Some of the observations described above have been used to explain

the behavior of larger castings and ingots. In Type A alloys the first

y precipitates in the primary grain boundaries and as the temperature

drops, intragranularly as well. All of these regions of y have different

orientations and the columnar-to-equiaxed transition is thus observable

27

in the primary (solidification) and secondary (solid state) microstructures

.

In Type B alloys, where y begins to solidify from the melt after initial

6 solidification and prior to completion of solidification, columnar y

crystals nucleate at high temperature and grow under steep temperature

gradients across the entire ingot cross section. With the high driving

force for y growth produced by the temperature gradient, changes in the

21 32 33 original 6 orientation present no barrier to that growth. ’ *

Mixtures of Type A and Type B behavior are frequently observed in the

same ingot. Once 6 has solidified, the y envelope is difficult to form

27

at high freezing rates. Therefore, Type A solidification might be

expected in the vicinity of the mold wall and Type B solidification in

the ingot interior. Variations in 6-ferrite level across an ingot cross

section must be rationalized in light of the fact that higher 6-ferrite

levels, associated with higher cooling rates, can be offset by rapid

homogenization in the finer solidification structures, which lowers the

20 39 6-ferrite content. 5 As the 6-ferrite levels in welds and castings

43 44 increase, the 6 phase becomes more continuous. ’ Spheroidization

Pf) of 6 can also take place during the 6->y transformation.

Measurement of 6-ferrite Content

There are a number of methods that have been employed to determine

the 6-ferrite levels in duplex austenitic stainless steels and most have

371
10 45

been reviewed recently. ’ The commonly employed methods are magnetic

measurements and quantitative metallography. Quantitative metal lography

is not as useful for measurements of the fine 6-ferrite structures of

welds, and extensive counting is necessary to compensate for the wide

1

6

intrinsic variations in 6-ferrite content within a given microstructure.

Several more exotic techniques for 6-ferrite measurement have been

1 5 attempted. Of these, integrated Mossbauer spectra and electrolytic

43 separation may prove particularly useful because they are not dependent

on the shape, size, orientation or distribution of the 6-ferrite.

RESULTS OF METALLOGRAPHIC INVESTIGATIONS OF CASTINGS AT NBS

The bulk of metallographic investigations on castings conducted at

46

NBS to date has been performed on CF8M material (ASTM A351 ) . These

alloys are a cast version of AISI 316 stainless steel. Alloy designation,

ferrite level, and chemical composition of the CF8M alloys are contained

in Table II. These castings were keel blocks, shown in Figure 7. All

keel blocks were annealed at approximately 1400 K.

The macrostructure of these alloy castings is typified by the

CF8M-9 alloy. The macrostructure, delineated by a light color etch,* is

shown in Figure 8 and is characterized by a region of columnar dendritic

grain growth perpendicul ar to the casting surface, followed by a transition

to an equiaxed dendritic growth form. The color etch delineates the

solidification structure by depositing a blue interference film on the

regions of the original 6 dendrites (high Cr) and a gold film on the

interdendritic regions (low Cr) . When viewed in color, the columnar

region to the left of the transition boundary possesses a blue cast

3 *C^lor Etch - A deposition etch composed of: 100 cc distilled water, 20

cc hydrochloric acid, and 2 g potassium metabisulphite.

372
whereas the equiaxed region to the right possesses a gold cast, indicative

of a greater volume of interdendritic liquid in the latter region. A

series of higher magnification micrographs depicts the nature of the

transition in solidification growth morphology (Figure 9a, b, and c).

Note that the dark gray areas correspond to the blue original 6-ferrite

dentrites and the light gray areas correspond to the gold interdendri ti

c

liquid regions. The region adjacent to the casting surface is composed

almost completely of 6-ferrite dendrites, aligned perpendicul ar to the

casting surface (Figure 9a). The microstructure in the transition

region is characterized by 6-ferrite dendrites that grow at angles to

the heat flow direction, effectively blocking columnar dendritic growth

in this direction (Figure 9b). The transition also appears to signal an

increase in the volume of interdendritic liquid in the as-solidified

structure. This correlates with the gold appearance of the central zone

of the keel block casting when viewed at lower magnification. The

structure in the central portion of the casting is characterized by

randomly oriented columnar dendrites as well as equiaxed dendrites

(Figure 9c). The finer network structure observed inside the original

6-ferrite dendrites is the residual 6-ferrite present at room temperature

subsequent to the S^y transformation.

When a heavier color etch is applied to the same casting section a

coarse columnar grain structure parallel to the direction of heat flow

is observed (Figure 10). The columnar grain structure appears to grow

across the previously described col umnar-to-equiaxed grain morphology

transition. A higher magnification macrograph of the columnar crystals

373
shows that the dendritic structure, as defined by the residual 6-ferrite,

possesses the same orientation across many of the columnar grain boundaries

(Figure 11). The columnar grain boundaries also appear faceted on a

microscal e.

A 10% oxalic acid electro-etch is employed to examine the residual

6-ferrite morphology since this etch preferential ly attacks the 6/y interphase

boundaries. Some characteristic features of the 6-ferrite morphology

are immediately apparent from observations of the CF8M-9 alloy (Figure 12).

On a given plane of polish the 6-ferrite appears composed of either

relatively long, filamentary particles (e.g., at A in Figure 12A) or

equiaxed cell-like networks (e.g., at B in Figure 12a). The 6-ferrite

at this volume fraction also appears to be composed of discrete particles

that are faceted in many instances. A facet is defined as a straight

segment of 6/y interphase boundary. There are many regions of apparently

fine, equiaxed 6 particles (e.g. at C in Figure 12a) and other regions

where 6/y interfaces are tangent to one another (e.g., at D in Figure

12a). The latter regions would seem to be locations of incipient spheroidization

of larger 6-ferrite particles. A less prevalent feature of the microstructure is the apparent twinning of the y matrix between the regions

of resiudal 6-ferrite (Figure 12b). The location and morphology of the

6- ferrite is obscured in these regions when the oxalic acid etch is

employed.

The faceted 6-ferrite particles also seem to be characterized by

definite angles that frequently occur between different facet variants

(Figure 13). The frequently occurring intervariant angles are 60° and

90°. Although no independent assessment was made of the crystal lographic

374
plane in the 6-ferrite parallel to the plane of polish, these intervariant

angles are consistent with the interplanar angles of different 110

6-ferrite planes ({110)6. This means that faceted 6/y interfaces could

be parallel to {110)6. Comparison of the residual 6-ferrite microstructure

with the original 6-ferrite dendrite morphology provides some independent

confirmation of this notion. A micrograph of a color etched region of

the CF8M-9 alloy illustrates this relationship for facet traces with a

60° angle between them (Figure 14). The primary dendrite axis has been

47 established parallel to a <100>6 easy growth direction by other investigators.

The measured angle between the primary dendrite axis and the facet trace

is 45° (see the schematic diagram in Figure 14). The assumption of the

facet plane parallel to {110)6 is then consistent with the measured

angle of 45° since this is one of the tabulated angles between {110)6

and {100)6 in a bcc crystal. X-ray and TEM determination of the crystallographic orientations of 6 and y are planned to confirm the accuracy of

this analysis.

The residual 6-ferrite distribution was employed to define the

orientation of the original 6-ferrite dendrite spines in Figure 11. As

shown in Figure 15, the 6-ferrite located at the dendrite spines is

actually composed of a network of 6-ferrite.

375
As the 6-ferrite content of the CF8M alloys is increased, the

6-ferrite microstructure on a given plane of polish becomes coarser

while still possessing many of the characteristics observed in the

CF8M-9 alloy, i.e., facets, points of interface tangency, networks

(Figure 16). A ION KOH electro-deposition etch is employed to deposit

a light blue to light brown film on the 6-ferrite. This provides greater

contrast between the 6 and y phases while eliminating the etching of y

twins. The coarser 6-ferrite microstructures are also characterized by

a greater apparent continuity. The 6/y interface morphology is quite

intricate and can generally be described as straight facets connecting

curved segments of interface. Closer scrutiny of the 6/y interfaces

indicates the presence of ledges on both faceted and curved segments

( Fi gure 17).

DISCUSSION

The metal lography of the CF8M alloys is consistent with the Type B

solidification sequence: L-d_+6->L+6+y->6-i-Y . In addition to residual

6-ferrite located at dendrite spines, this sequence is distinguished

metallographically by a secondary structure, due to solid state transformation

,

in which the col umnar-to-equiaxed grain morphology transition of solidification is not observed (Figures 8 and 10). Figure 8 shows the primary

(solidification) structure with its col umnar-to-equiaxed transition;

376
Figure 10 shows the secondary (6->y) structure composed of completely

columnar crystals that bend parallel to the direction of heat flow. The

fact that dendrites are observed to grow across the columnar grain

boundaries of the secondary structure also supports the assertion that

the columnar y structure is a secondary structure produced by the 6-*y

transformation (Figure 11). In the case of a solidification sequence

where a y envelope begins to solidify from the melt, columnar y crystals

can nucleate heterogeneous! y at high temperature and grow with the aid

of very steep temperature gradients. Once nucleated under these conditions,

a columnar y crystal will grow right across misorientations between

6-ferrite dendrites and parallel to the path of maximum heat flow (Figure

10). These metallographic observations are also consistent with respect

to the various Ni- and Cr-equival ent ranges specified for primary 6

alloys in Table I

.

The larger volume fraction of primary 6-ferrite near the casting

surface is partly attributable to the higher cooling rate in this region

and the difficulty in forming the y envelope at high solidification

27

rates. The interruption of 6 solidification by y solidification has

been associated with the columnar-to-equiaxed transition, in which large

32

pockets of liquid exist between settling 6 crystals (Figure 9).

The level of residual 6-ferrite does not appear to vary significantly

with distance from the casting surface and neither does the dendrite arm

spacing (Figure 9). The latter is a monitor of local solidification

48 time and determines the time available for a diffusion controlled

transformation to occur. Finer spacings might be expected to promote

377
the 6->y transformation and thereby decrease the level of residual 6-ferrite.

Therefore, the constancy of both residual 6-ferrite content and 6

dendrite arm spacing implies that the cooling rate does not vary sufficiently

within the casting to affect these parameters significantly, although it

still might affect the interruption of 6 solidification by y solidification.

A general model for the 6->y transformation is proposed to rationalize

the observed microstructural features of the residual 6-ferrite in

primary 6 alloys. The model assumes that the 6+y transformation is

thermally activated and diffusion controlled. A nucleus of y forms

either at solidification subgrain boundaries or on the y envelope formed

by direct y solidification from the melt. These alternative modes of

solid state y nucleation would affect the nucleation rate and distribution

of y nucleii but not the basic morphological evolution. The shape of a

nucleus and the implications of faceted nucleation and growth have been

49-52

considered generally and with specific regard to proeutectoid a.

Analogous nucleation and growth is assumed to occur during the 6->y

transformation. The occurence of a facet on a nucleus is associated

with formation of a low-energy interface between parent and product

phase. The low energy interface, furthermore, is associated with a high

degree of lattice matching (coherency) across the interface and a crystallographic

orientation relationship between the abutting phases. The segment or

segments of low energy interface minimize the total surface energy of the

nucleus and thereby minimize the activation barrier for nucleation.

373
This type of nucleus has been used in models of Widmanstatten-type

proeutectoid a formation (saw teeth and side plates).

49-52

In this kind

of proeutectoid y-*y transformation, the nucleus and precipitate faceting

41

is associated with a Kurdjumov-Sachs (K-S) type of orientation relationship.

A K-S orientation relationship, between bcc and fee phases, involves a

parallelism between planes and directions of densest atomic packing in

the product and parent phases. For the K-S orientation relationship the

following parallelisms exist:

(Hlb // { 11 0}a

<11 0>y // <1 1 1 >a

Faceting is definitely an aspect of 6/y interface morphology in the

castings of the present investigation. Furthermore, optical metal lographic

evidence suggests that the facet plane in the 6-ferrite parent phase is

a {11 0 } 6 (Figures 13 and 14). This is a plane of densest packing in a

bcc phase and is highly suggestive of a K-S orientation relationship

existing between y and the 6 it is growing into. As mentioned in the

previous section. X-ray and TEM will be employed to confirm whether or

not a K-S orientation relationship exists between 6 and y in these

alloys. This orientation relationship has been observed for so-called

'lathy* 6.

20

Once y nucleus formation has been considered, it remains to demonstrate

the y growth process necessary for the observed evolution of 6-ferrite

microstructure. Assuming a distribution of y nucleii and facet morphologies,

growth is accomplished mainly through the movement of the curved, incoherent

segments of 6/y interface. These are high energy and high mobility

379
boundaries. The faceted, coherent segments possess low mobility since

49-52

atom jumping across the interface is energetically infeasible.

Movement of a faceted boundary perpendicul ar to itself can be envisioned

49-52

as proceeding by the movement of ledges or steps parallel to the boundary.

These steps are actually incoherent boundary segments that possess high

mobility. The presence of ledges in a macroscopi cally curved interface

moving along specific crystallographic planes can also provide a means

for maintaining an orientation relationship between the parent and

product phases separated by the interface. Therefore, a faceted interface

is not necessary for observation of a crystal lographic orientation

relationship between parent and product phases. Ledges have been observed

on both faceted and curved 6/y interfaces in this investigation (Figure 17).

The features of the residual 6-ferrite microstructure can then be viewed

as evolving through the rapid movement of incoherent segments of 6/y interface

with coherent segments being dragged along behind. The coherent segments

move outwards as well, by a ledge mechanism, but more slowly. Points of

tangency between transformation fronts after impingement are commonplace

and accelerate spheroidization processes. A schematic diagram showing

the evolution of the residual 6-ferrite morphology is shown in Figure

18. It is important to note that for both the model and the actual

6-ferrite microstructures observed in this study, observations on a

given plane of polish can be deceiving. These types of observation

might lead one to the erroneous conclusion that discrete 6-ferrite

particles are formed, even at high volume fraction. In fact, observations

380
43

in the author's laboratory and by other investigators suggest that a

continuous 6-phase exists at contents greater than about 7 v/o. Therefore,

although spheroidization of 6-ferrite certainly does take place at high

temperature, the "discrete" particles and points of interfacial tangency

observed may be peculiar to only that plane of polish (Figure 18). The

characteristic filamentary 6-and cel 1-1 ike 6-ferrite morphologies observed

(Figure 16B) could then be considered as two views of a continuous

6-phase (Figure 18). The ferrite network can be more equiaxed if impingement of incoherent segments occurs more uniformly, and this will certainly

occur if the solidification structure is finer and the y nucleation rate

higher. These conditions might be expected with welding and other higher

cooling rate processes. Observations of the residual 6-ferrite microstructures

in primary 6 weld metals show them to be more equiaxed, i.e., to exhibit a higher

volume of cell-like than filamentary 6-ferrite. The cell-like networks still

possess faceted 6/y interfaces (Figure 19). The similarities between

residual 6-ferrite microstructures in welds and castings indicate potential

similarities in their solidification and 6-»y transformation behavior.

RECOMMENDATIONS FOR FURTHER WORK

The following are recommendations for further work:

1) TEM, selected area electron diffraction and X-ray diffraction studies

to determine the fine structure of 6/y interfaces and the crystallographic orientation relationships that exist between 6 and y.

2) Scanning transmission electron microscopy (STEM) studies to determine

the concentration gradients of Ni and Cr that exist within the microstructures of duplex austenitic stainless steels. Knowledge of the

381
Ni and Cr concentration profiles across 6/y interphase boundaries

may assist in elucidating the nature of the 6+y transformation.

3) Quantification of the proposed model for 6->y transformation to

confirm its ability to predict observed volume fractions of residual

6-ferrite within the constraints of normally encountered cooling

rates.

332
CONCLUSIONS

1. The CF8M castings in this study are primary 5 solidification alloys

with the following transformation sequence: L-*L+6-*L+6+y->6+y

2. The primary structure of these alloys is characterized by a columnarto-equiaxed grain morphology transition in the casting interior.

3. The secondary structure, owing to the 6->y transformation, is characterized

by columnar crystals that extend through the entire casting cross

section, bending in the direction of heat flow. The secondary

structure completely masks the presence of the primary structure,

as noted in previous investigations.

4. The residual 5-ferrite microstructure is generally characterized by

curved segments of 6/y interface connected by straight or faceted

segments of 6/y interface. The 6-ferrite becomes coarser and

appears to possess greater continuity as the 6-ferrite content

increases.

5. Frequently occurring interfacet angles of 60° and 90° are consistent

with a { 1 1 0 } 6 crystallographic plane parallel to the facet plane

and the establishment of a Kurdjumov-Sachs orientation relationship

between 6 and y across the faceted 6/y interface.

6. The observation of ledges on curved and faceted 6/y interfaces

demonstrates the possibility of interface movement via a ledge

mechanism that could maintain an orientation relationship across

any type of 6/y interface*

7. The features of the residual 6-ferrite morphology in this type of

alloy can be rationalized with the aid of a transformation model

analogous to that proposed for the proeutectoid y->a transformation.

383
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386
27. J. M. McTighe and J. Beech: "The Structure of As-Cast Austenitic

Stainless Steels," J . Res , SCRATA , No. 37, (June 1977), p. 43.

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32. H. Fredriksson: "The Solidification Sequence in an 18-8 Stainless

Steel, Investigated by Directional Solidification," Met. Trans

.

,

Vol. 3, (11), (1972), p. 2989.

33. A. Hultgren: " Crystal 1 ization and Segregation Phenomena in 1.10%

Carbon Steel Ingots of Smaller Sizes," J. I .S .

I

. , Vol. 70 (1929),

pp. 69-113.

34. P. Lichtenegger and R. Bloch: "Color Etching of High Alloy Steels,"

Prak. Metal!

.

Vol . 12, (1975), p. 567.

35. J. C. Lippold: PhD Thesis, Rensselaer Polytechnic Institute, Troy,

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307
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6-ferrite into Austenite in Continuously Cast Slabs of 18-8 Stainless Steel," Tetsu-to-Haqane , Vol. 65, (8), (1979), pp. 1176-1185.

40. J. B. Newkirk: "General Theory, Mechanism and Kinetics" in Preci pitation

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41. G. Kurdjumov and G. Sachs: "Crystal lographic Orientation Relationship

between a-and y -Fe," Ann. Phys. , Vol. 64 (1 930), p. 325.

42. T. B. Massalski: "Massive Transformation, Chapter 10" in Phase

Transformations , ASM Seminar, American Society for Metals, Metals

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Delta Ferrite in Austenitic Stainless Steel Weldments by an Electrochemical Technique," Wei d. J . Res . Suppl

.

, Vol. 58, (12), (1 979),

p. 375-s.

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Ferrite Contents," Wei d. J . Res . Suppl . , Vol. 57, (6), (1 979), p.

167-s.

388
45. R. B. Gunia and G. A. Ratz: "The Measurement of Delta Ferrite in

Austenitic Stainless Steel," Weld. Res. Counc. Bull. 132, (1968).

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Fabrication of the Cold Bore Tube for a Large Superconducting CFFF

MHD Magnet", paper presented at the 1980 Superconducting MHD Magnet

Design Conference, MIT, Cambridge, Massachusetts, March 26-27,

1980.

47. J. W. Christian: "Solidification and Melting," Chapter 14 in

The Theory of Transformations in Metals and Alloys , Pergamon Press,

New York (1965), pp. 562-566.

48. M. C. Flemings: "Solidification Processing," Met. Trans

.

, Vol . 5,

(2), (1974), p. 121.

49. W. C. Johnson, C. L. White, P. E. Marth, P. K. Ruf, S. M. Tuominen,

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Met, Trans. A , Vol. 6A, (4), (1975), p. 911.

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Cementite Reactions" in Decomposition of Austenite by Diffusional

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Growth of Precipitate Crystals, Chapter 8", in Phase Transformations ,

ASM Seminar, American Society for Metals, Metals Park, Ohio, (October

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Meeting, Niagara Falls, New York (1976).

389
Table I. Classification of Solidification Modes in

Austenitic Stainless Steels

Investi gation Classi fication Classification Composition Appl ication

References 16, 27 L+L+6+6 Type A Cr = 18.2-24.2

Nigq = 8.6-11.2

L+L+6->L +6 +y^-6 +y Type B Cr = 18.1-21.7

Ni^ q = 10.1-13.6

eq

Castings

L+L +y_>L +y +<5 -*6 +y Type C Cr = 18.9-20.3

Ni

q = 13.9-26.6

eq

L->L+y^y Type D

References 19-22 L->L+y^y Type A Cr /Ni s 1 .48

eq eq

Welds

L+L+6+6 Type C Cr /Ni > 1.95

eq eq

L^-L+6->L+6+y^6+y Type B 1 .48 s Cr /Ni s 1 .95

eq eq

Reference 29 L->-L +

6

->L +6 +y^y + <S $ < 0 Castings

L^-L+y + <S^Y +<S $ = 0 * Ni

eq

- °- 75Cr

eq

L->L+y~>y $ > 0 + 0.257
Table II. Alloy Designation, Chemical Composition,

6-ferrite Level

Chemical Composition (w/o)

Alloy Designation C Mn Si Cr Ni Mo v/o 6-ferriti

CF8M-9 .02 1.06 1.10 19.33 11.67 2.14 9

CF8M-1

.05 .83 1.25 19.50 12.20 2.69 14

CF8M-24 .07 .52 1.32 21 .48 9.48 3.18 24

391
List of Figures

Figure 1 (a) Schaeffler Diagram (b) DeLong Diagram.

Figure 2 Schoefer Diagram.

Figure 3 Constant Fe pseudobinary phase diagram.

Figure 4 Schematic diagrams of primary 6-ferrite solidification

modes.

Figure 5 Schematic diagram of solute distribution after normal

freezi ng

.

Figure 6 Schematic diagram of solute distribution acr^s a subgrain,

following primary 6 solidification of a weld

Figure 7 Schematic diagram of keel block casting.

Figure 8 Keel block cross section showing col umnar-to-equiaxed

grain morphology transition in solidification structure,

1 ight color etch.

.

Figure 9(a) Keel block cross section in columnar grain region near

edge, light color etch.

Figure 9(b) Keel block cross section near col umnar-to-equiaxed grain

morphology transition in solidification structure, light

color etch.

Figure 9(c) Keel block cross section in equiaxed grain region, light

color etch.

Figure 10 Keel block cross section showing columnar y crystals

formed during 6-»y phase transformation, heavy color etch.

Figure 11 Columnar y crystal structure in relation to solidification

structure delineated by residual 6-ferrite, heavy color

etch.

Figure 12 Residual 6-ferrite morphology in CF8M-9 alloy (a) filamentary

and cell-like 6-ferrite, (b) twinning in y, oxalic acid

electro-etch.

Figure 13 Residual 6-ferrite faceting, oxalic acid electro-etch.

Figure 14 Residual 6-ferrite in relation to 6-ferrite dendritic

solidification structure, light color etch.

Figure 15 Residual 6-ferrite networks at 6-ferrite dendrite spines,

light color etch.

392
Figure 16 Residual 6-ferrite morphology (a) CF8M-14, (b) CF8M-24,

ION KOH electro- deposition etch.

Figure 17 Residual 6-ferrite in CF8M-24 alloy showing ledge morphology,

ION KOH electro-deposition etch.

Figure 18 Schematic diagram of evolution of residual 6-ferrite

morphology via 6^y phase transformation

.

Figure 19 Residual 6-ferrite in weld metal; ~3v/o 6-ferrite,

316L GMA, oxalic acid electro-etch.

393
Ni

Equivalent

=

Ni

Equivalent

=

(%Ni+30%N+30%x%C+0.5x%Mn)

(%Ni+30%C+0.5%Mn)

Figure 1 (a) Schaeffler Diagram (b) DeLong Diagram.

394
RATIO

CHROMIUM

FACTOR

TO

NICKEL

FACTOR

VOLUME % FERRITE

Figure 2 Schoefer Diagram.

395
Primary y

Solidification

Primary 8

Solidification

Figure 3 Constant Fe pseudobinary phase diagram .

396
8 + y 8 + L

Final

Structure

(a) Type A ( L-*L+6-*6->-6 +y)

Envelope

L+B+y -L+8

(b) Type B ( L->-L+6-*-L+6+y->-6+y)

Figure 4 Schematic diagrams of primary 6-ferrite solidification modes.

397
COMPOSITION

DISTANCE

Figure 5 Schematic diagram of solute distribution

after normal freezing.

398
Boundary Core

Figure 6 Schematic diagram of solute distribution across

a subgrain, following primary 6 solidification

of a wel

d

^

' .

399
Figure 7 Schematic diagram of keel block casting.

400
Transition

Transi tion

/

3 cm

Figure 8 Keel block cross section showing columnar-toequiaxed grain morphology transition in

solidification structure, light color etch.

401
Columnar Grain Growth Dir

500 jam

Figure 9(a) Keel block cross section in columnar grain

region near edge, light color etch.

402
Columnar Region Transition Equiaxed Region

I

500 um

Figure 9(b) Keel block cross section near columnar-toequiaxed grain morphology transition

in solidification structure, light

color etch.

403
500 pm

Figure 9(c) Keel block cross section in equiaxed

grain region, light color etch.

404
Figure 10

2 cm

Keel block cross section showing columnar

y crystals formed during 6-*y phase transformation,

heavy color etch.

405
20 cm

Figure 11 Columnar y crystal structure in relation

to solidification structure delineated

by residual 6-ferrite, heavy color etch.

406
250 urn

Residual 6-ferrite morphology in CF8M-9

alloy (a) filamentary and cel 1-1 ike

6-ferrite, (b) twinning in y, oxalic

acid electro-etch.

Figure 12
Fi gure

Y'

\ 4

9. ' Ci=C>

' » « d

o T/V :j

-

rj

j • ; V

• ~'V ' <3x

•

' ^4/ •

V

p •..

& t 0 N>1

A

° r» $

J ><f

>

j z *.

X;

to <P'* *

"! ?

•’* '>>

>

. - -X * FX.

a >.

. ,

' j

' :X*

f'.'; $ V' f

' '• , & r

~ o , , e^-^jT ” r -*

.

< . . \ . ' - r

*

, X . '• -' T ft ' • 7 < x- * ;

- ",

x '> > „• ' . / ft ---X ^ ' /

' "V '-.V

\ > - • '

.. v ft V. ft A • c.e ' i •>>•<> f ^

<? i • v

-

. 1. ' 2* V

’

„*=* Xo (J'.

'','-'' * F

' 'ft ->

* v "N* ft! ft>‘

- ‘ -

'

*>• -''X 77

i l "X A

.r-x>

* <s /X

' c£>* - i# •> X :, O- >'•/,, .

’ >

trX -

' , , / xx>h

x> • o "x.

' ftft '-ftX

v- *

•

X

''v.

’• s, x " -X"

tt° ' > v, C-ft ^ <5 \>-X; - Xa1 -, Xx- ^

i> / 2 ft

' x .

.. : -

'

- ; > - — x

X - « ' \

- A X „ <7

/.

250 liffl

13 Residual 6-ferrite faceting,

oxalic acid electro-etch.

408

A

<f7
250 pm

Facet variants

of residual

8-ferrite

Secondary

dendrite arm

<100>g growth

direction of 8

dendrite

(primary dendrite arm)

Figure 14 Residual 6-ferrite in relation

to 6-ferrite dendritic solidification

structure, light color etch.

409
Figure 15 Residual 6-ferrite networks at 6-ferrite dendrite

spines, light color etch.

410
V

(a)

(b)

250 urn

Figure 16 Residual 6-ferrite morphology (a)

CF8M-1 4 , (b) CF8M-24 , ION KOH electrodeposition etch.

411
(a)

31 y m

Figure 17 Residual o-ferrite in CF8M-24 alloy

showing ledge morphology, ION KOFI electro-deposition etch.

412
Ajepunoq

mej6qns

.

<50

>*

m

c

o

E

03

413

Figure

18

Schematic

diagram

of

evolution

of

residual

6-ferrite

morphology

via

6->y

phase

trans

formation

.
, . • ' -IS j , vc V V> / - .

-

. .v^. v

r f * .. >v*r -

*• ,-•*

r<* • * ..." ~ ' X'- * . >. •

‘ S> » ' C JP V * . f +-' — ; < v . *

. -\J " - x V'V.

- ,

• -• • rf v * 'v -,' * A. • ..

c*.

4 r

' •-;- ">:->• AV

r.r* •

.

- 1 . ^ ,

.

j*r * - C V. • v- >A-

» * «s i

/ x

> , *< V-> " ’ i

c¥*‘~ *

w- % -

*

t v

y

. ^

f jV '*> x • A" -v$ r *

x * % ^ *

.

# V

' ' *

" >*X >

* ^

j

•' ^ *„ % ;

' sK -

</ . s

f »*..*’

*-"' '•

f

’ * » '* -'A

*

* % r #/* • A

C > • ^ •» > > t 'n*

. r ^ *

1 25ym

Figure 19 Residual 6-ferrite in weld metal; ~3v/o

6-ferrite, 316L GMA, oxalic acid electro-etch.

414
DEFORMATION AND FRACTURE OF STAINLESS STEEL CASTINGS AND

WELDMENTS AT 4 K

T. A. Whipple and E. L. Brown

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

A number of stainless steel weldments and castings that were deformed at 4 K have been examined with optical and electron microscopy.

The purpose of this investigation was to assess qualitatively the effects

of residual delta-ferrite on the deformation and fracture mechanisms of

these materials at cryogenic temperatures.

Weldments and castings have very similar microstructures , but the

ferrite in castings is much coarser as a result of slower solidification

and cooling rates. This wide variation in the scale of the structure

provides a good opportunity for the study of structure and property

relationships in these materials.

It has been found that deformation and fracture are greatly influenced by the ferrite morphology. The study of the coarser structure in

castings has allowed more direct observation of fracture, which has led

to some preliminary conclusions concerning fracture mechanisms. Cleavage

fractures of ferrite particles have been observed in some castings and

in an electraslag weldment. It appears that these fractures, although

not forming a continuous crack path, assist in the void formation in the

ductile austenite matrix. Another observed crack propagation mechanism

is that of void nucleation near the austenite/ferrite interface. Although

interface separation has not been observed, the inhomogeneous deformation

near the interface appears to cause premature void nucleation and growth.

The occurrence of one of the above mechanisms rather than the other

depends on a number of factors, including ferrite morphology and the

properties of the two phases. The morphology and properties, in turn,

depend on both the alloy composition and solidification conditions.

415

INTRODUCTION

The structural components for most superconducting magnets are

being made of stainless steels. Because of the large structures, high

stresses and low operating temperature, welded construction is the most

attractive fabrication method. Since the structures are being made of

thick plate, fracture toughness is an important consideration. It has

been found that stainless steel weldments generally have a much lower

1 2 fracture toughness than the stainless steel base metal. ’

To date, three factors have been established that have an effect on

the cryogenic fracture toughness of stainless steel weldments: the

1 3 delta-ferrite content of the weldment, * the nitrogen concentration in

4 2 the weld, and the fusion zone grain size of fully austenitic weldments.

Nevertheless, it still is not possible to predict the toughness of a

stainless steel weldment on the basis of these factors. Apparently

there are factors beyond those stated above that cause the observed wide

2 5 6 variations in fracture toughness. *

* This paper contains the results

of the initial stage of an investigation that is intended to define more

precisely the role of ferrite and to identify other factors that have an

influence on the cryogenic toughness of stainless steel weldments and

casti ngs

.

In many structural applications it would be advantageous to use

stainless steel castings to facilitate fabrication and to eliminate the

need for welding. There are presently very few 4 K mechanical property

data on castings, but a few initial tests have shown a wide variation in

5 fracture toughness similar to weldments. This is not suprising

417
because weldments are actually rapidly solidified castings. Stainless

steel castings also generally contain a certain percentage of del ta-ferri te

,

although on a much coarser scale than that of weldments, owing to the

slower solidification of castings.

This paper presents the results of a metal lographic and fractographic

investigation of deformation and fracture mechanisms in duplex austenite/

ferrite structures at 4 K,

MATERIALS

The materials examined in this study are stainless steel weldments

and castings that have been tensile and fracture toughness tested at

4 K. These materials did not come from a controlled series, however,

they all fall into the AISI type 304 or 316 stainless steel composition

ranges. These variations in composition, as well as variations in the

casting or welding process, are appropriate since the intent of this

study was to investigate the possible deformation and fracture mechanisms

in duplex austenite/ferrite structures at 4 K and not just the mechanism

for a specific material. The available chemical composition and mechanical

property data for the materials used in this study are presented in

Tables 1 and 2.

PROCEDURES

For this investigation a number of sectioning, metal 1 ographic and

fractographic techniques were employed. Each technique allows the

observation of different phenomena. Macroscopic examination, optical

microscopy, and scanning electron microscopy (SEM) were used for making

observations.

418
Sectioning

Many of the figures presented in this paper are taken from samples

that were sectioned perpendicular to the fracture plane of tensile and

compact tension specimens. This sectioning method allows the observation of crack propagation paths and deformation around the crack tip and

provides insight into the crack propagation mechanisms.

Metal 1 ographic Specimen Preparation

For the most part metal 1 ographic specimens were prepared by using

standard mechanical grinding and polishing techniques, but in a few

cases electrolytic polishing was used.

A 10% oxalic acid electrolitic etch was used as a general purpose

etch to reveal ferrite morphology and grain boundaries. This etch can

also reveal strain-induced martensite if the etching time is increased;

however, if the specimen contains ferrite, this procedure leads to considerable surface relief and pitting, which make metal 1 ography difficult.

A 10 N KOH electrolytic etch was used to reveal ferrite structure

only. This is a deposition etch that colors the ferrite light blue to

brown. The main attribute of this etchant is that it does not remove

any material; therefore, it will not round the edges of cracks or produce

pitting and other etching artifacts. The KOH etch is sometimes used

prior to oxalic acid etching because it tends to passivate the ferrite

and allow deeper oxalic acid etching of the austenite with less overall

damage.

A solution of 50 m£ HC1 , 125 m£ H^O and 2g Na^SO^ was used to color

etch stainless steel weldments and castings. This etchant is useful in

419
revealing the original solidification pattern after the solid-state

ferri te-to-austeni te transformation has occurred.

A magnetic etching procedure was used to reveal the magnetic phases,

ferrite, and bcc martensite. For this procedure a suspension of fine

magnetic particles is spread on the surface of the specimen and the

specimen is placed in a magnetic field; the magnetic particles will then

migrate to the magnetic phases. This etch was used mainly in conjunction

with the KOH etch to reveal bcc martensite which is an indication of

plastic deformation in the austenite matrix.

DISCUSSION

Low Temperature Deformation

It is generally true that bcc materials, such as delta-ferrite, go

through a ductil e-to-brittl e transition at low temperatures, whereas fee

materials, such as austenite, do not. This is due to the difficulty of

operating non close-packed slip systems at low temperatures. The temperature at which this transition occurs in steels is influenced mainly

by alloy content and microstructure; a higher nickel content and a

smaller grain size lead to a lower transition temperature. Another

factor to consider is the increase in strength with a reduction in

temperature; bcc materials have a much more dramatic increase in yield

strength than fee materials at low temperatures

.

These general observations lead to some expectations for the behavior

of duplex stainless steels. At higher temperatures , both austenite and

ferrite are ductile and have similar strengths; in this case the amount

420
or morphology of the ferrite is not expected to significantly affect the

strength or toughness of the material. At temperatures below the transition

temperature of the ferrite, it becomes a strong, brittle second phase.

Factors that influence the inhomogeneity of deformation in the ductile

austenite matrix will influence the strength and toughness of the material. This means that the amount, morphology and properties of the

ferrite can all have significant influences. Observations that support

these hypotheses have been made by Read et al.

1

Figures 1 and 2 are micrographs of 316L stainless steel weld metal

tensile specimens that were tested at room temperature and 4 K, respectively. Note that at room temperature (Figure 1) the ferrite is highly

elongated in the direction of the tensile axis and at 4 K (Figure 2) the

ferrite maintains its original dendritic morphology and shows no indication of elongating with the matrix.

During deformation at low temperatures, austenite can go through a

strain induced martensitic transformation. The degree of transformation

and the critical strain at which it occurs are a function of the alloy

content. Although the stability of the austenite matrix will quantitatively effect the strength and toughness of duplex austenitic stainless steel castings and weldments, it is only of interest in this

qualitative discussion in that it provides a method of observing deformation in the austenite.

Figures 3 and 4 are micrographs of areas near the crack tips in two

CF8* stainless steel castings tested at 4 K. The intent of these micro-

*CF8 and CF8M are ASTM A351 grades of cast stainless steel and are analogous to AISI 304 and 316, respectively.

421
graphs is to show the inhomogeneity in deformation caused by the presence of ferrite. Figure 3 is from a 0% ferrite casting; the straight

lines near the crack tip are strain-induced martensite. Note that the

lines are straight and fairly uniform, only changing direction at grain

boundaries. This can be contrasted with Figure 4, which is a similar

region in a 14.5% ferrite casting. In this case, the martensite formation is extremely nonuniform and seems to be concentrated in association

with the ferrite morphology. Another example of the effect of the

presence of ferrite on deformation patterns in the austenite matrix is

shown in Figures 5 and 6. In Figure 5, bending of martensite bands near

the ferrite particles can be seen in a CF8M* casting with 9% ferrite,

indicating that the ferrite is nondeforming and the matrix is flowing

around it. Figure 6 is from a 316L electroslag weldment; this micrograph shows the reinforcement of the austenite matrix by the ferrite in

a low strain area. Note that there is little martensite formation

between the parallel ferrite plates and more extensive martensite

formation where the matrix is not as well reinforced.

All of the above evidence lends support to the observations cited

at the beginning of this discussion: (1) at low temperatures ferrite is

a very strong nondeforming phase and (2) the ferrite content and morphology

are important influences on how the ductile austenite matrix deforms.

Low Temperature Fracture

As a result of the above discussion on 1 ow-temperature deformation,

it appears appropriate to classify duplex austenitic stainless steel, at

*CF8 and CF8M are ASTM A351 grades of cast stainless steel and are analogous to AISI 304 and 316, respecti vely

.

422
low temperatures, as a ductile matrix composite with hard, brittle

particles or discontinuous fibers. This is not an exact classification,

since the ferrite is very seldom in the form of spherical particles and

can be either continuous or discontinuous depending on solidification

and cooling conditions, but this classification does provide a good

starting point for the discussion of fracture mechanisms. There are

three possible mechanisms for the fracture of ductile matrix/harddiscontinuous fiber composites.^

1. Brittle fracture of the fibers.

2. Matrix/fiber interface separation.

3. Fiber pull-out.

Brittle fracture of the fibers occurs when the local stress in the

fiber exeeds its fracture strength. The local stress in the fiber is

not necessarily the same as the macroscopic stress on the sample, but

can be considerably higher owing to inhomogeneous plastic deformation in

the matrix. Interface separation is an important consideration with

artifical composites and with metal -matrix/particulate composites;

however, the interface between austenite and ferrite is a very strong

one and no evidence of interfacial separation has been observed. Fiber

pull-out is sometimes associated with interface separation, but for the

purpose of this discussion it means a ductile failure of the matrix near

the matrix/fiber interface. This occurs when the fiber does not fracture,

but the highly inhomogeneous deformation of the matrix near the fiber

results in void nucleation at this location. The occurrence of one of

these mechanisms over the others depends on fibre strength, content and

morphology.

423
In this study evidence was found for both fiber fracture and fiber

pull-out. Until this investigation there had been no direct observation

of the brittle fracture of ferrite in weldments and castings at 4 K.

Although this had long been assumed to be the cause of the reduction in

toughness of stainless steel weldments with increasing ferrite content,

direct observation was difficult because of the extremely small size of

the ferrite in weldments. In this investigation observations were made

on both castings and weldments made with high heat input processes. The

much slower solidification and cooling rates of these materials results

in a coarser structure, which is more easily observed.

With these materials extensive evidence was found of brittle fracture

of ferrite in weldments and castings deformed at 4 K. Figure 7 is a

micrograph of a 316L electroslag weldment taken near the fracture surface

of a tensile specimen. In this specimen, numerous brittle fractures of

ferrite particles are evident. The fact that the fractures are at 90°

to the tensile axis lends further support to the conclusion that they

are brittle cleavage fractures. Additional evidence for cleavage fracture

of ferrite is shown in Figure 8 (a graph from the 14.5% ferrite CF8

casting). This micrograph shows consistant angular relationships between

the fractures, indicating that fracture is occurring on a specific set

of crystallographic planes, i.e., cleavage planes.

Figure 9 and 10 are crack profiles for two stainless steel castings. Figure 9 is a CF8M casting that contains 9% ferrite and has a

very high toughness at 4 K, similar to that of wrought base metals.

Figure 10 is a CF8 casting containing 14.5% ferrite; this casting has a

low toughness comparable to that of a weldment with a similar ferrite

content.

424
In both cases the crack path is associated with the ferrite morphology,

but Figure 10 has a much more brittle appearance and significant cracking

not associated with the ultimate crack path. Also, it can be seen in

Figure 4 (a higher magnification of the crack tip) that there is significant

ferrite cracking in advance of the crack tip, which was not observed in

the CF8M casting of Figure 9. A definitive explanation of the difference

in these two castings cannot be offered at this time; however, two

factors appear to have contributed the difference in the amount of

ferrite cracking. First, the ferrite morphology is different; the

tougher CF8M casting was a sand casting with a slower solidification and

cooling rate than the centrifugal ly cast CF8 casting. This resulted in

a coarser and discontinuous ferrite morphology, which provided less

restriction to deformation in the austenite matrix than the finer more

continuous ferrite in the CF8 casting. The second contributing factor

may be the toughness of the ferrite phase: CF8M has a higher nickel

content than CF8, and it is well known that nickel increases the cryogenic

toughness of ferritic materials.

Weldments differ from castings mainly in that the solidification

and cooling rates are much greater. This results in a more directional

solidification pattern and a much finer structure. The greater cooling

rate results in less diffusion during the ferrite- to-austenite transformation

causing the ferrite to be more closely associated with the original

dendritic solidification pattern. Therefore, the residual ferrite of

weldments is more likely to be continuous. When fracture is associated

with cracking of the residual ferrite, having continuous ferrite can provide

425
a preferred crack path and lower toughness than occurs for a material of

the same ferrite content in a discontinuous morphology.

An example of this preferred crack path is shown in Figure 11 in

the crack profile of a fracture toughness specimen from a 316L electroslag

weldment. An electroslag weldment experiences a cooling rate greater

than that of a casting, but much less than that of other welding processes.

This results in a relatively coarse but continuous ferrite morphology.

Figure 11(a) shows that the dendritic growth direction determined the

crack propagation direction even though the stress state was such that

the crack should propagate straight. The closer view. Figure 11(b),

shows the crack was propagating down dendrite cores and secondary dendrite

arms. The crack periodically moved to an adjacent dendrite owing to the

stress state. Figure 12(a) is a micrograph of the crack tip of this

weldment, showing that the crack is indeed propagating in a brittle

manner through the ferrite. Although this fracture does have a very

brittle appearance, the measured fracture toughness was above average

for a weldment of this ferrite content. Figure 12(b) shows the same

area using the magnetic etch to reveal strain-induced martensite. Even

though the ferrite fractures in a brittle manner, there was still considerable

plastic deformation in the austenite matrix. This matrix deformation

and the fact that energy was consumed in the crack propagating off-line

accounts for the relatively high toughness of this weldment. Further

evidence that the ferrite in this weldment fractured in a brittle manner

is provided by the SEM fractographs in Figure 13. These fractographs

demonstrate the mixture of brittle ferrite fracture and ductile matrix

failure.

426
Behavior similar to that described above can also be seen in the

rapidly solidified 308L flux-cored metal arc (FCMA) weldment shown in

Figures 14 and 15. Because of the rapid solidification, the ferrite is

extremely fine and it was not possible to observe brittle fracture of

the ferrite directly. Figure 14 shows that the crack propagation again

followed the dendrite growth direction. The closer view of Figure 15

shows that the crack path was determined by the local ferrite morphology,

which varies considerably from the bottom to the top of a weld pass.

As mentioned previously, two possible discontinuous fiber composite

fracture mechanisms are of interest: fiber fracture and fiber pull-out.

The evidence already presented has shown that fiber fracture does occur

in coarse duplex structures. In the case of weldments with much finer

structures, limited evidence has been found for both fracture mechanisms. Figure 16 is a stereo fractograph of the fracture surface of the

308L FCMA weldment described previously. The area indicated by the

arrow appears to be a subsurface ferrite dendrite, where the inhomogenous deformation nucleated a ductile failure near the ferrite/austenite

interface. Composite fracture theory predicts that, for a discontinuous

fiber composite, there is a critical aspect ratio (fiber length-todiameter ratio) at which a transition from fiber fracture to fiber

pull-out fracture mechanisms occurs; a high aspect ratio favors fiber

fracture. The aspect ratio of a fiber is dependent on orientation to

the tensile axis: the same fiber will have a high aspect ratio if it is

aligned with the tensile axis and a low one if it is perpendicular to

the tensile axis. In Figure 16 it is apparent that the dendrite was

perpendicular to the tensile axis; therefore fiber pull-out occurred.

427
Figure 17 shows a case where the fibers were apparently parallel to the

tensile axis and fiber fracture occurred.

SUMMARY

Samples of stainless steel castings and weldments that had been

deformed and fractured at 4 K were examined. The thrust of the investigation

was to determine the mechanisms by which residual delta-ferrite affects

the deformation and fracture of duplex austenite/ ferri te structures at

cryogenic temperatures. Evidence was presented that indicates that at

low temperatures the ferrite acts as a hard, brittle second phase particle.

The manner in which these particles affect the mechanical properties of

the material depends on the amount, properties and morphology of the

second-phase particles. It was observed that the ferrite reinforces the

austenite matrix and causes the deformation to be inhomogeneous. Evidence

was found for two fracture mechanisms: (1) brittle fracture of the

ferrite mixed with ductile failure of the matrix between ferrite fractures

and (2) ductile failure in the matrix near the austeni te/ ferrite interface,

which was caused by the highly inhomogeneous deformation in this region.

The occurrence of one of these mechanisms over the other depends on the

morphology and properties of the ferrite.

Future work in this area will concentrate on obtaining a more

quantitative description of the role of ferrite in the deformation and

fracture of duplex austeni te/ ferrite materials at liquid helium temperatures.

428
REFERENCES

1. D. T. Read, H. I. McHenry, P. A. Steinmeyer, and R. D. Thomas,

"Metallurgical Factors Affecting the Toughness of 31 6L SMA Weldments at Cryogenic Temperatures," Weld. J. Res. Suppl . 59 , 1 04s113s (1980).

2. T. A. Whipple, H. I. McHenry, and D. T. Read, "Fracture Behavior of

Ferrite-Free Stainless Steel Welds in Liquid Helium," to appear in

Wei d. J. Res. Suppl . 60 (4), 1 981.

3. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic Stainless Steel Weld Metals: Part I - Role of Ferrite,"

Weld. J. Res. Suppl. 57, 325s-333s (1978).

4. E. R. Szumachowski and H. F. Reid, "Cryogenic Toughness of SMA

Austenitic Stainless Steel Weld Metals: Part II - Role of Nitrogen,"

Weld. J. Res. Suppl. 58 , 34s-44s (1979).

5. T. A. Whipple and H. I. McHenry, "Evaluation of Weldments and

Castings for Liquid Helium Service," in Materials Studies for Magnetic

Fusion Energy Applications at Low Temperatures-I

V

, NBSIR 81-1645,

pp. 273-287 (1981 ).

6. T. A. Whipple and D. J. Kotecki, "Weld Process Study for 316L

Stainless Steel Weld Metal for Liquid Helium Service," in Material

s

Studies for Magnetic Fusion Energy Application at Low TemperaturesIV., NBSIR 81- 1645, pp. 303-321 (1981 ).

7. A. Kelly, "Particle and Fibre Reinforcement," in Strengthening

Methods in Crystals , Halsted Press Division, John Wiley & Sons,

New York, pp. 433-484 (1971).

429
LIST OF TABLES

Table 1. Process and chemical composition data on stainless steel

weldments and castings.

Table 2. Mechanical properties of stainless steel weldments and castings.

430
Table

1.

Process

and

chemical

composition

data

for

stainless

steel

weldments

and

castings.

cr» LO

z: O O

J 0 1 1 O 0 0

O

C\J CO 00

Q- O 0 0 0 1 1 O 0 0 0

1 8

00 CXJ

CNJ r— r— 1 1 OO O O O O 1 8 O O O O

+>

r\ O 00 O

c O 1

0 1— CO 1 1 o « • •

1 1 •r— O 0 CXI r—

4-> •i—COO

ClEO •r— O co O

O • • • • 8 B

j r— 00 1— r— 1 1 r— 1

r— 1—

03O

•r—E

CD -C CO o i- O CO CO CO O • • • • 8 8 CO cr> cr> CO 1 1 r— 1—

C\J CXI O

•r— r^ CO 1— co 1 1 OO • • • •

1 1

1 r— 0

CT» CO LO C 0 1— 0 1 | 21 • • • • 8 8

1 O 1 '

CO LO

C\J 0 (XI OJ O 0 0 0 O 1 1

1 1 0 0 0 O

r* CD

4-> LO 0 LO CNJ •r— • • • • • S- 0 ^3* cr> <3- 00

S- r—

<D

Ll. *0

1 CD r— 1— CO

05 03 03 "O “O

CO cn CO CO co r— O r- cd a 3 C 3 C c CO CD CD £- £- M— -r- M— •r— O CO ^ O 03 E •r- 4-> •r- 4-> 4-> “O S- a 1 S- S- CO S- CO CO -*-> r— <D 0 1 1—

o 4-> fO -*-> 03 “O rd O CD E S- X 03

L±_ c a C O c 0 CD ^ _Q n3 3

CD CD 03 r— 3 »— CD O O 00 UJ OO Lc- E

>>O 2: -J —J -J r— CO 00 00 CO 00 00 r— Ll_ Ll. u_ r— O 0

c O O 0 CO CO CO

431
Table

2.

Mechanical

properties

of

stainless

steel

weldments

and

castings

at

4K.

|C

« > -

-

«. —>. * .. *

-—x *r— 1— oo CM 1— CM •o CO O r— LO CM -X CO i— CM i— j

o

[—

1

— — —

* Lh

03 1— LO 00 CD CM CD

CL. CO CM r^. 00 CO r-^ s: CO 1 CM 1 1

d

o *

•r- fC

4-> CD O S- 3 c ^ 1 I i CM

“O

CD d

q: °r—

1 1 8 CM CM CM

do

°r—

4-3 03 CD CO CO CD 00

CD^-S do

CO 1 CM

I

UJ

*> »r— i— CD CO r^ CM LD

CD -d CO CO r— CD CD CD

4-> 4-> 1— CM i

1

i

i—

03 CD E c

"

x ' " v—" 1 " ' "

•r- CD o CM 00 LO 00 P i- (3 LO i

CM LD CO LO

«— 4-> Q_ CM LO CO OO f

OO DOOS i

•r- o 00 CD O CO

* CO CO CD O CO CM

-d ^ r— 1— i

“0 4-3 r— CD

v—" 'x v—

CD d

“r— CD 03 CO o LO o CO LO

>- S— Q- r— r^ CD 0 LO

CO

•=d" LO CD CD 00

CD

4-3 LO o LO r-. CM

•r— • • • • • o CD 00

,

CD

Li_ ~a

i—

(D

i

i— CD

03 CD 03 CD CD 03 *o

CD d CD d d i— ~a CD O

=3 • r— 3 •«“ •r— to CD S- s-

<4— 4-> 4- 4-3 4-> O -O CD O 03 E •r— CO •r- to CO S- r— S- o

S- S- 03 S- 03 03 4-3 CD CD 1 1—

O 4-> O 4-3 O “O o o ^ E X 03

Ll_ d d d CD -O 3 4-> CD CD 03 r— 3 r- CD

CJ> O oo LU OO Li- £=

>>o _J —

J

I f— CO CO CO CD CD 00 r— U- u_ Li_ i

,— o

<c o o O CO CO CO

432
LIST OF FIGURES

Figure 1 Micrograph of a 316L SA weld metal tensile specimen

deformed at room temperature, taken near the fracture

surface and showing the elongation of the ferrite. KOH

etch.

Figure 2 Micrograph of a 31 6L SA weld metal tensile deformed at

4 K, showing that the ferrite does not deform at low

temperatures. Oxalic acid etch.

Figure 3 Area near the crack tip of a 0% ferrite CF8 casting.

Oxalic acid etch.

Figure 4 Area near the crack tip of a 14.5% ferrite CF8 casting.

KOH and oxalic acid etch.

Figure 5 Micrograph of a 9% ferrite CF8M casting deformed at 4 K,

showing the deformation of the austenite matrix around a

ferrite particle. Color etch.

Figure 6 Micrograph of a 316L electroslag weldment deformed at

4 K, showing the rei nforcement of the austenite matrix by

the ferrite particles. KOH and magnetic etch.

Figure 7 Area near the fracture surface of a tensile

specimen from a 316L electroslag weldment. Oxalic acid

etch.

Figure 8 Area near the crack of a 14.5% ferrite CF8 casting. KOH

and oxalic acid etch.

Figure 9 Crack profile of a 9% ferrite CF8M casting. Color etch.

Figure 10 Crack profile of a 14.5% ferrite CF8 casting. KOH and

oxalic acid etch.

Figure 11 Crack profile of a 316L electroslag weldment. Oxalic

acid etch.

433
Figure 12

Figure 13

Figure 14

Figure 15

Figure 16

Figure 17

Micrographs of the crack tip of a 316L electroslag weldment,

a) KOH etch, b) KOH and magnetic etch.

SEM fractographs of the fracture surface of a 31 6L

electroslag weldment.

Macroscopic crack profile of a 308L FCMA weldment. Oxalic

acid etch.

Micrograph of the crack tip region of a 308L FCMA weldment.

Oxalic acid etch.

SEM stereo fractograph of a 308L FCMA weldment.

SEM fractograph of a 308L FCMA weldment.

434
TENSILE AXIS

50 ym

Figure 1 Micrograph of a 316L SA weld metal tensile

specimen deformed at room temperature, taken

near the fracture surface and showing the

elongation of the ferrite. KOH etch.

435
I

TENSILE

AXIS

I

200 y m

Figure 2 Micrograph of a 31 6L SA weld metal tensile

specimen deformed at 4 K, showing that the

ferrite does not deform at low temperatures

.

Oxalic acid etch.

436
437

CF8

casting.

Oxalic

acid

etch.
•MI ‘

• _ M\r: .

^ ^ - > • V ~

' 4H ! “t'Z

f .. ? %<

P- o

438

Figure

4

Area

near

the

crack

tip

of

a

14.5%

ferrite

CF8

casting.

KOH

and

oxalic

acid

etch.
50 urn

Figure 5 Micrograph of a 9% ferrite CF8M casting

deformed at 4 K, showing the deformation of

the austenite matrix around a ferrite particle.

Color etch.

439
100 n m

Figure 6 Micrograph of a 316L electroslag weldment

deformed at 4 K, showing the rei nforcement of

the austenite matrix by the ferrite particles.

KOH and magnetic etch.

440
Figure 7 Area near the fracture surface of a tensile

specimen from a 31 6L electroslag weldment.

Oxal ic acid etch.

441
50 ym

Figure 8 Area near the crack of a 14.5% ferrite CF8

casting. KOH and oxalic acid etch.

442
443

Figure

9

Crack

profile

of

a9%

ferrite

CF8M

casting.

Color

etch.
444

Figure

10

Crack

profile

of

a

14.5 %

ferrite

CF8

casting

KOH

and

oxalic

acid

etch.
(b) 400 ym

gure 1

1

Crack profile of a 31 6L electroslag weldment.

Oxal ic acid etch

.

445
(b)

50 ym

e Figure 12

*

'oqraphs of the crack tip of a 316L

\ 4- N l/nu Cstrh hi ctroslag weldment

magnetic etch.

a) KOH etch, b) KOH

446
Figure 13 SEM fractographs of the fracture surface of a 31 6L electroslag weldment.

447
(b) 1 mm

Fiqure 14 Macroscopic crack profile of a 308L FCMA

weldment. Oxalic acid etch.

I

LOADING

DIRECTION

I

448
449

Micrograph

of

the

crack

tip

region

of

a

308L

FCMA

weldment.

Oxalic

acid

etch.
Fi gure 1 6 SEM stereo fractograph of a 308L FCMA

weldment

.

450
Figure 17 SEM fractograph of a 308L FCMA weldment.

451

*

EVALUATION OF WELD FILLERS TO JOIN NITROGEN STRENGTHENED STAINLESS STEELS

FOR CRYOGENIC APPLICATIONS

(FUSION AND MHD ENERGY)

R. H. Espy

Armco Inc., Research and Technology

Middletown, Ohio

ABSTRACT

Tests are reported of austenitic stainless steel weldments that use

filler materials manufactured outside the United States.

453

Evaluation of Weld Fillers to Join Nitrogen Strengthened

Austenitic Stainless Steels for Cryogenic Applications

(Fusion and MHD Energy)

R. H. Espy - Armco Research & Technology

US/USSR Technical Exchange Program

on Welding Stainless Steels

BACKGROUND INFORMATION

During the exchange visits of the Russian and United States Cryogenic

Welding Groups both in the USA and Russia, information on weld fillers

for joining the Nitrogen Strengthened Austenitic Stainless Steels

was presented. While the weld compositions studied in the U.S.A. had

acceptable strengths and toughnesses, there was some concern about the

presence of micro fissure defects that were almost always present.

The concern was mainly that of the effect of defects under cyclic

loading conditions.

Initially the cooperative effort between the U.S. and USSR was to

exchange sample weldments having defects typical of those encountered

in welding fully austenitic stainless steels. During our first ex- change it became evident that the defects we were talking about were

metallurgical in nature (micro fissure) and the defects the Russian

delegation was referring to were of the workmanship type like incomplete penetration, unfused or unwet areas, cold laps, entrapped slag,

etc. Actually, we did not consider workmanship defects as being

important because of our nondestructive test and repair practices

which are very common and effective throughout the USA where high

quality weldments are desired. On the other hand, the Russian group

did not consider metallurgical defects like micro fissures because

they did not occur with the weld fillers they were using. In general

one might say that Russian weld metallurgy in this area appeared more

advanced than in the U.S., but on the other hand that U.S. manufacturing technology appeared more advanced than the Russian manufacturing capability. The Russian approach of applying large numbers

of people to technological problems appears to be the reason for this.

In Russia's Paton Welding Institute alone there were, in 1979, 34

engineers and 17 technicians assigned to welding the stainless steels.

As a first step in resolving this major difference we inquired con- cerning the typical composition of the Russian weld filler. This was received along with a coil of bare wire for our laboratory

evaluation. Just prior to our exposure to Russian technology in

the exchange program, we at Armco decided to examine two foreign

made weld fillers touted to be good for weldments used for

455
2 .

Cryogenics. One designated as Grinox 41 (stick electrodes) and

Grinox 5 (bare wire) is made by Messer. Griesheim in Germany

(distributed by MG Weld Products - Menomonee Falls, Wisconsin)

The second designated as P6 (stick electrodes & bare wire) is

made by Avesta in Sweden (distributed by A. Johnson & Co . - Lionville, Pennsylvania). The results of our work with these

two wires are included with those of the Russian wire in this

report.

COMPOSITION OF WIRES S WELD DEPOSITS

Analyses of the filler wires and weld deposits (for covered

electrodes) showed them to have the following compositions.

A typical composition of fully austenitic weld fillers studied

in the U.S. is shown for comparison. .

C Mn P S

Russian

Bare Wire .027 6.7 .01 .00 3

German

Grinox 41

Weld Deposit

.17 9.9 .012 .004

Swedish

Avesta P6

Weld Deposit

.03 4.9 .015 .001

U.S.

Armco

Weld Deposit

.05 9.0 .007 .009

PROCEDURE USED IN PREPARATION OF

Si Cr Ni Mo W N

.09 19.45 16.1 2.8 2.1 .23

CO 0 15.6 12.4 - 2.3 -

.25 18.5 17.5 2.6 - .17

.22 18.0 16.2 .18

ALL-WELD TEST SPECIMENS

The weld geometry employed was as specified in AWS 5.4-79 except the

base plate was 1" thick, the root opening was 3/4" and the buttered

overlay was deposited as one layer using 5/32" <f> Type 309 Stainless

covered electrodes . The parameters for welding with the bare wire

were as shown.

Weld Process - Automatic Gas Metal Arc (Pulse)

Filler Wire - .062" <|) X Coil

Shield Gas - Argon + 2% O

2 at 30 CFH

Weld Current - 160 amperes

Volts - Average, 19.5

Peak, 66.0

Cycles - 120

Weld Travel Speed - 15 IPM

Stringer Bead Practice Used (No Oscillation)

Weld Interpass Temperature RT/300°F

456
3 .

The parameters for welding with covered electrodes were as shown.

Weld Process - Shielded Metal Arc (Manual)

Filler - 5/32" <f> Weld Current - 100 amperes

Weld Voltage - DCRP 24 volts

Stringer Bead Practice Used

Weld Interpass Temperature RT/300°F

The .505" <f> tensile and standard Charpy V-notch specimens were

sectioned from the top center of each deposit. The notch in each

Charpy specimen was oriented top to bottom of the weld deposit.

TEST RESULTS

All of the items welded in a satisfactory manner with the processes

employed

.

The room temperature and -320°F tensile and Charpy data are shown

below. Included again for comparison are data obtained with the

U.S . weld filler.

Tests at Room Temperature

Tensile Toughness

Defects De fects

UTS .2% YS % El on Fracture CVN on Fracture

ksi ksi 2" % RA Face '# Face

Russian 104 73 41 56 None 114 1 ppt.

German

Grinox 41

95 73 35 65 Some o 11nU)

Separations

116 None

Swedish

Avesta P6

95 66 41 63 None 95 None

U.S. Armco 93 71 25 26 1S% (2) 44 25%

M.C. M.C.

Tests at -320°F

Russian 186 133 19 26 5% (2) 68 3 small ppts

M.C.

German

Grinox 41

173 118 31 31 C „(!) Some 11

Separations

62 None

Swedish

Avesta P6

179 123 40 40 None 37 None

U.S. Armco 162 134 10 12 10% (2) 13 20%

M.C. M.C.

(1) Appear to be heat affected area grain boundary separations

parallel to direction of welding.

(2) M.C. (Micro Condition) appear to be areas of grain boundaryseparations and weakness associated with the freezing pattern.

457
4 .

Me tallographic examination of the weld structures to determine

reasons for soundness or conversely lack of soundness , showed

the molybdenum and tungsten bearing weld metals to contain coring

or a phase that reacted sufficiently different to electrolytic

etches so as to stand out significantly. These weld deposits

were also free of grain boundary breaks in the high temperature

heat affected areas of the multipass welds. The sensitized or lower temperature affected areas of the German Grinox weld

contained many grain boundary precipitates of carbides. The

tungsten and molybdenum free welds (U.S. - Armco) showed numerous grain boundary breaks in the high temperature heat affected

zones. These structural conditions and their affect on tensile

and impact tests are shown in Figures 1 to 4.

X-ray diffraction examination of the areas in the tungsten and

molybdenum containing welds that showed a differential reaction

to the etchants showed all areas to have the same crystal

structures. The SEM analytical efforts were not successful in

showing molybdenum and tungsten segregation, the elements

believed responsible for the etching phenomena. Microprobe

efforts with a Phillips AMI/3 Electron Microprobe to identify

element separation were not successful. A new electron microprobe unit, Cameca CAMEBAS at Armco, with which it is believed

that proper identification of the highlighted areas can be made

will be in service shortly. The results from this effort will

be reported separately when available.

DISCUSSION

The use of tungsten and molybdenum together as in the Russian

wire or separately as in the German and Swedish electrodes is

effective in controlling the high temperature hot short con- dition found to occur in the welds made with the U.S. - Armco

electrodes

.

All of the welds showed increased yield strengths at cryogenic

temperatures confirming the fact that both N&C influence yield

strengths as temperatures are lowered. The German Grinox

electrode contained residual nitrogen (low) with C near .20%

while all other Russian, Swedish and U.S. fillers contained N

at about .20% with C at a low level.

While all of tungsten and molybdenum containing welds were

acceptably sound and exhibited good properties at all temperatures there was a trend in the high carbon German alloy towards

a tensile fracture pattern that reflected the presence of grain

boundary carbides in lower temperature or sensitized areas of

the multipass welds. While the Russian filler showed low indications of porosity and micro condition, both it and the Swedish

alloy appeared overall quite good.

458
5 .

CONCLUSIONS AND COMMENTS

The use of tungsten and molybdenum in ferrite free stainless weld

metal is effective in controlling the grain boundary breaks that

have been found to occur in high N containing welds.

Normally ferrite in high nitrogen stainless welds is an effective

means for producing sound defect-free welds, but ferrite when

present has an adverse effect on toughness at cryogenic temperatures

when associated with the high strength imparted by nitrogen.

While all ferrite-free stainless steel welds are prone to form some

grain boundary breaks, the use of nitrogen for strength caused

significant increases. The data reported shows that both tungsten

and molybdenum are effective in controlling cracking with the processes used.

The German Grinox 41 (stick electrode) and Grinox S (bare wire) and

Swedish Avesta P6 (stick electrode and bare wire) are commercially

available for any added weld process work that may be desired.

Armco is currently making an effort to produce the Russian wire so as to make it also available for additional weld process test work.

Overall the US/USSR exchange program has been beneficial in that

the higher nitrogen stronger stainlesses like 304 LN (.15 N) and

NITRONIC 40 (.30 N) can be considered for weldments at the liquid

helium temperatures associated with fusion and MHD energy projects.

The initial objective of determining the effect of weld defects

still remains to be addressed and should be considered in future

exchanges. While the use of tungsten and molybdenum were

effective in controlling the nitrogen induced grain boundary defect,

there were indications that a small amount of the defect still

occurred and this may vary with the weld processes employed.

It is suggested that the program to determine the effect of

defects be continued in future exchange visits.

R.H.Espy

10/20/80 459
Mag 8X

Fracture Face of CVN

Tested at -320°F Mag 1.5X

Tensile Tested at

Room Temperature

Etchant NaOH Mag 500X

FIGURE 1 : Impact, Tensile and Microstructure

for Weld Deposited Using Russian

Wire

C Mn P S Si Cr Ni

027 6.7 .01 .003 .09 19.5 16.1

N Mo W

.23 2.8 2.1

Toughness at -320°F

CVN Defects

68'# 3 small ppts

Tensile at Room Temperature

UTS .2% YS %E1

ksi ksi 2" %RA Defects

104 73 41 56 None

Coring or Phase in Ferrite-Free

Austenite

460
Fracture Face of CVN

Tested at -320°F Mag 1 . 5X

Tensile Tested at

Room Temperature

FIGURE 2 : Impact, Tensile and Microstructure

for Weld Deposited Using German

Grinox 41 Electrodes

c Mn P S Si Cr Ni W

17 9.9 .012 .004 .80 15.6 12.4 2.3

o Toughness at -320 F

CVN Defects

62 ' # None

Tensile at Room Temperature

UTS

ksi

95

.2% YS

ksi

73

%El

2" %RA Defects

35 65 Some 11

Separations

Etchant NaOH Mag 500X

Coring or 2nd Phase in

Ferrite-Free Austenite

461
Tensile Tested at

Room Temperature

FIGURE 3 : Impact, Tensile and Microstructure

for Weld Deposited Using Swedish

Avesta P6 Electrodes

C

.03

Mn

4.9

P

.015

£

.001

Si

.25

Cr

18.5

Ni

17.5

N

.17

Mo

2.6

Toughness at -320 F

CVN Defects

37'# None

UTS

ksi

95

Tensile at Room Temperature

Defects

.2% YS

ksi

%E1

2

"

66 41

%RA

63 None

Coring or 2nd Phase in

Ferrite-Free Austenite

462
Mag 8X

Fracture Face of CVN

Tested at -320°F

Mag 1 . 5X

Tensile Tested at

Room Temperature

FIGURE 4 : Impact, Tensile and Microstructure

for Weld Deposited Using U.S. - Armco Electrodes

c Mn P S Si Cr Ni N

05 9.0 .007 .009 .22 18 16.2 .18

Toughness at -320°F

CVN Defects

13'# 20% Micro Condition

UTS

ksi

93

Tensile at Room Temperature

.2% YS

ksi

71

%E1

2” %RA

25 26

Defects

15% Micro

Condition

Etchant Chromic Mag 500X

Grain Boundary Defects ir

Ferrite-Free Weld Metal

463

HOHMETALUCS

465

NONMETALLICS FOR MAGNET SYSTEMS

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

An introduction to the nonmetal lies program is presented,

discussed are: standard grades, radiation resistance, materials

characteri zation , and technology transfer. Plans for FY 81

are presented.

Topics

467

The nonmetal lies program has continued to respond to the needs of

magnetic fusion energy (MFE) designers and fabricators by working with

industry to provide commercial sources of standard grades of insulating

laminates, emphasizing those grades that are most widely used in critical

parts of superconducting magnets. The Annual Report for FY 79 (NBSIR

80-1627) reported the establishment of specifications for cryogenic

grades of glass-fabric reinforced bisphenol A epoxy industrial laminates

identified as G-10CR and G-11CR. The initial mechanical, elastic,

electrical, and thermal property characterization data were reported

to 4 K. This work has continued during FY 80, concentrating on refining

the specification to achieve a balance between minimum performance

variability at 4 K and minimum cost to the purchaser.

It was recognized at the outset that radiation resistance of

the insulating laminates would be an important performance requirement.

The G-11CR product was incorporated as an initial effort to achieve

better radiation resistance than that of the more common and lower cost

\ G-10 type laminates. Available data had indicated that the aromatic

amine cure of the G-ll product should be beneficial. However, tests

performed in the Oak Ridge reactor failed to confirm this expectation.

Because this left the G-11CR, a fairly expensive material, with no overall

performance advantage over the less expensive G-10CR, the G-11CR

product was dropped from the list of standard laminate grades.

G-10CR is, therefore, the sole standard cryogenic grade at this

time. Reception of this product by the cryogenic industry has been good,

and it is presently being manufactured by three major U.S. laminating firms.

469
As originally conceived, the standard grades were to provide the basis

for systematic material development in response to fabricators' needs,

in addition to providing baseline materials having predictable cryogenic

properties. Thus, a variant of the basic G-l OCR product i ncorporati ng

a lightweight glass cloth has been developed by one of the manufacturers

.

This variant, designated G-10CR-L, was required to increase the number

of reinforcement layers in a thin wall thermal standoff for a Lawrence

Livermore Laboratory magnet design.

Magnet fabricators also require a uniaxial glass-reinforced laminate

product and would find it advantageous to have available a glass-matreinforced epoxy laminate having proven performance at 4 K. The uniaxial

product would primarily serve in the fabrication of thermal standoff support

straps, and the mat product would be, in many instances, a lower cost

substitute for G-10CR, having the added advantage of isotropic properties

in the plane of the laminate. Substantial progress has been made in

working with industry to establish standard specifications for such materials.

Consideration has also been given to the possibility of establishing a lamina'

coding system that would permit the designer to specify the least expensive

commercial material required to meet his needs without having to be intimately

familiar with all of the parameters affecting insulating and structural

laminate performance. This is long-range effort which, if successful, would

encompass both standard cryogenic grades and the proprietary products of

various manufacturers

.

These considerations are further expanded in a publication "Standardization

Of Composites for Cryogenic Service" reproduced in this section of the report.

470
An example of industry interaction and participation is given by the

discussions between NBS representatives and two industrial laminators

during the NBS-DoE Vail Workshop that addressed refinement of the G-10CR

specification. A copy of the Vail presentation by J. R. Benzinger

entitled "The Manufacture of Radiation Resistant Laminates" is reproduced

with the permission of Mr. Benzinger to provide background on the capability

and limitations of industry in meeting the performance requirements of

insulators for MFE magnets.

The materials characterization portion of the nonmetallics effort

was continued in FY 80 with additional studies of the effects of cryogenic

temperatures on the thermal conductivity of G-10CR and G-11CR laminates.

An effort to measure conductivity of the epoxy matrices separately was

only partly successful owing to inhomogeneity in the bulk resin. Such

data would be of value in establishing models for predicting the conductivity

of laminates of differing construction.

A glass-fabric-reinforced industrial laminate of European manufacture

that had previously been evaluated at CERN was included in the NBS

thermal conductivity program to provide a comparison of results between

the two laboratories . For reasons not yet clarified, the NBS measurements

indicated about a 45% higher conductivity at room temperature than did

the experiment at CERN. It appears that the differences were due to

experimental discrepancies between the laboratories rather than to

specimen variabilities. Work to define the source of the differences is

conti nui ng

.

471
The effect of temperature on the failure mode of insulating laminates

must be understood if accurate predictions of the life-cycle performance

of magnets are to be forthcoming. The document "Fracture Analysis of

Compression Damage in Glass Fabric/Epoxy, G- 1 OCR , 295 K to 4 K" included

in this report reflects an initial effort in this direction. Compression

failure was selected for study, since such stresses predominate in

operating magnets.

The high anisotropy of composite laminates creates problems in

measuring elastic constants by conventional strain gage methods. The

effort to develop dynamic methods of obtaining such measurements has

therefore been continued during FY 80. Such methods are attractive

because the small specimen size facilitates the evaluation of anisotropic

effects. But the necessary conditions for the dynamic moduli to be the

technical equivalent of engineering moduli obtained by conventional

techniques has not been established. The progress made during the past

year in clarifying the equivalency is documented in the contribution

entitled "Static and Dynamic Moduli of Two Fiber-Reinforced Composites".

The review paper entitled "Cryogenic Properties of Fi 1 amentary-Rei nforced

Composites: An Update" that appeared in the FY 79 report has been revised

in part and submitted for publication to the journal Cryogenics . The

major change was addition of a section on the radiation effects on

composites, a copy of which is included in the present report.

Two miscellaneous contributions were made to the literature during

FY80. A letter to the editor of the journal Composi tes comments on a

contribution by K. Kadotani relating to the effect of cryogenic temperatures

472
on the performance of composite materials. Work at NBS has indicated

that the authors' conclusions could lead to serious misjudgments in

materials selection if they are misinterpreted. We also replied to an

invitation from the American Society for Metals to contribute to their

forthcoming issue "Technology Forecast '81'. Copies of both these

documents appear in this report.

We anticipate that the nonmetallics program for FY 81 will continue

to emphasize the areas of materials standardization, materials research

and characterization, and technology transfer. The standardization

effort will pursue the introduction of standard cryogneic grades of

glass mat and uniaxial glass-reinforced epoxy laminates. The characterization program will screen the performance of a G-10CR product from a

third producer as well as that of a lightweight glass variant of this

product. The performance of a polyimide-matrix variant of G-10CR will

also be screened, because this laminate has shown substantially better

radiation resistance than the epoxy grades. The research efforts will

involve a cooperative program with industry to produce a cycloaliphatic

epoxy resin variant of the basic bisphenol A G-10CR for future studies

of radiation resistance. In-house research will attempt to resolve the

noted discrepancy in laminate thermal conductivity and will continue to

assess the validity of the dynamic method for determination of laminate

elastic constants. Work will also begin on a system for precise measurement

of viscoelastic effects in laminates as a function of temperature.

473

NONMETALLIC MATERIALS

STANDARDIZATION

475

STANDARDIZING NONMETALLIC COMPOSITE MATERIALS

FOR CRYOGENIC APPLICATIONS

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

The designer of superconducting magnet systems is confronted with a

difficult task in specifying the nonmetallic insulating and structural

support materials required for proper system functioning. He finds it

necessary to select from a profusion of products identified by trade

names, most of which were developed for other than cryogenic use and for

which reliable cryogenic performance data are nonexistent. This paper

describes the ongoing efforts by NBS and industry to provide a commercial supply of standardized cryogenic grades of the most widely used

insulating laminates having wel 1 -characterized cryogenic performance.

These laminates will also provide the basis for a systematic materials

development program in response to user needs. The elements required

for a more generalized laminate coding system i ncorporating all materials

of interest to magnet designers are also considered.

477

INTRODUCTION

The technology of composite materials is presently at the state

that metals technology was before being systematized by the establishment of industry standards. It is difficult to imagine working in

metals technology in the United States without the ability to associate

performance with designations such as 304L for stainless steel or

2024-T4 for aluminum. It would be equally difficult for individuals in

other countries to perform effectively without similar national metal

and alloy coding systems that transcend proprietary designations, but

this is the situation presently confronting those working in the composite materials area, in the United States and elsewhere.

Up to the present time, the problem has not been of major concern

because the consumption of high-performance composites has been relatively low. However, consumption is expected to increase greatly as

escalating energy costs place premiums on lightweight, fuel -effici ent

transportation systems and the development of alternative energy sources

It will become increasingly difficult to meet the future needs of composite technology with proprietary systems or with laminates fabricated

from the few formulations published in the literature.

The need for composite materials standards is therefore common to

many industries. As reflected in the forum chaired by Evans at the

1979 ICMC/CEC conference,^ the need of the cryogenic industry is particularly acute. This industry is already at the point where many of

the nonmetal lie materials needs of magnetic fusion energy (MFE), magneto

hydrodynamic (MHD) , and rotating cryogenic machinery systems cannot be

adequately met by existing commercial products.

479
A primary problem is the inability to associate reliably the cryogenic performance required for diverse applications with the profusion

of products identified by trade names, most of which were developed for

other purposes. This paper outlines the current thinking of the author

on this overall problem and describes the efforts currently underway to

meet the immediate industrial cryogenic requirements. The author hopes

that these considerations may be combined with similar considerations

based on European experience to contribute to the eventual establishment

of workable international standards for nonmetal lie materials.

The author will concentrate on nonmetallic composite materials,

reflecting the current need for such materials in the fabrication of the

cryogenic portions of new energy systems. It is recognized, however,

that unreinforced polymers pose a similar standardization problem.

IMMEDIATE CONCERNS

The immediate task is to assist designers and fabricators in the

selection of dependable insulating and structural nonmetallic laminates

for MFE and MHD systems presently being built in the United States. The

high costs of these systems and the substantial performance demands made

on the laminates require that the selected materials be well characterized at cryogenic temperatures and the variability in performance be

minimized. Because the present systems will provide experience that can

be extrapolated to more sophisticated, larger systems, the selected

materials must be commercially available over a period of years. Furthermore, they must lend themselves to a systematic materials development program in response to industry needs. These boundary conditions

480
imply that control over the constitutive elements, manufacturing processes, and performance characterization will be necessary.

Current construction will necessarily draw on state of the art,

commercially available nonmetallic laminate materials developed for

other than cryogenic service. In the past, many proprietary products

have been used. However, this approach is inefficient for demanding

applications since it requires that each group responsible for a design

conduct an expensive in-house materials characterization program. It

must also be assumed that the selected proprietary product will not be

substantially altered in response to market forces outside the cryogenic

area. The alternative of purchasing laminates to specified cryogenic

performance criteria either requires the manufacturers to develop cryogenic qual i ty-control procedures or assumes that room-temperature data

can be confidently extrapolated into the cryogenic range. Unfortunately, the latter concept is rudimentary for most conventional properties and nonexistent for special -performance requirements under irradiation at cryogenic temperatures.

A third alternative is to establish uniform industry specifications

for manufacturing a series of wel 1 -characteri zed nonmetallic laminates

having the greatest application to cryogenic technology. This approach

recognizes that many of the inexpensive commercial laminates developed

for room-temperature applications are useful cryogenic materials, whose

primary deficiencies are the lack of a cryogenic data base and an

excessive variability when used in a temperature region for which they

were not intended. Establishment and characterization of uniform specification laminates addresses these deficiencies and provides designers

with materials having greatly improved cryogenic reliability at minimum

481
cost. Furthermore, the uniform specifications assure that the critical

material factors affecting performance will be controlled by the cryogenic industry, providing the basis for a systematic materials development program that reflects the needs of the industry.

We have been following this approach in working with the United

States laminating industry to develop a series of standard, highpressure industrial laminates for cryogenic service. Both the industry

and the coordinating standards group, the National Electrical Manufacturers Association (NEMA) have been most cooperative. Thus far, two

grades of glass-fabric-reinforced epoxy laminates have been established,

characterized, and placed in commercial production. These materials are

designated G-10CR and G-11CR to distinguish them from the conventional

NEMA G-10 and G-ll products. Performance data for these cryogenic

2 grades have been published in the literature. Acceptance by the cryogenic industry in the United States has been gratifying. Work is continuing at NBS in cooperation with industry to develop additional

standard specifications for other laminate systems of interest to cryogenics, possibly including random-mat and uniaxial gl ass-rei nforced

epoxy-matrix laminates.

These developments are meeting the immediate needs of industry

while providing the basis for systematic alteration in the performance

of standard materials in response to developing industrial needs. The

concept of the overall program is outlined by the flow chart of Fig. 1.

Built-in iteration steps provide the open-ended type of program required

to be responsive to continually developing needs. The sequence depicted

for development of the CR grades of G-10 and G-ll illustrates the progress up to the present time. The characterization data indicate that

these products will perform very well in most applications.

482
Studies at the Oak Ridge National Laboratory, however, have shown

that the mechanical properties of these laminates are severely degraded

by 4-K neutron and gamma irradiation at fluences comparable to those

3 4

expected in a functioning MFE reactor. ’ This deficiency is being

addressed by a subprogram to develop laminates having improved radiation

resistance. The first step will be to determine if radiation resistance

will be significantly improved by replacing the epoxy matrix of the

G-10CR laminate with a polyimide matrix.

It appears to the author that such a program would also entail an

iterative cycle like the one illustrated in Fig. 2. Because the cryogenic radiation resistance of a polymer is likely to be a strong func5 tion of the exact molecular structure, an essential part of the cycle

is a correlation of molecular-level damage with radiation level and

property change. Only in this way will the required scientific basis be

established. The program envisions the development of methods for

preliminary screening of 4-K radiation resistance in the laboratory

prior to confirmatory testing in a reactor. It is evident that considerable new ground will have to be broken in solving the radiation

problems in nonmetal lie materials.

A major deficiency in the approach thus far discussed is the absence of an industry standards group having the responsibility for

coordinating, updating, and disseminating information on developed

standards. The NEMA organization is performing the disseminating

function for the CR-grade laminates through their Industrial Laminates

Subcommittee, but NEMA interest is confined to high-pressure , bulk

laminates. An organization responsive to the broader interests of the

composite materials industry is required.

483
LONG-RANGE STANDARDIZATION NEEDS

The above approach lacks the comprehensiveness required to assure

orderly overall composite materials development and implementation in

the long run. The author agrees with Evans that the most effective way

of alleviating the present chaotic state of nonmetallic materials technology would be to introduce a numerical classification system based

upon key elements of a material that affect performance.^ Here, the

cryogenic industry has much in common with other industries exploiting

composite technology.

There seems no a priori reason why such an approach would not be

feasible for composite laminates, when one recognizes that the significant engineering properties of a laminate at any temperature are largely

defined by relatively few parameters. For example, the intrinsic mechanical, elastic, thermal, and electrical properties of a laminate

should be defined within relatively narrow limits by a numerical classification system that establishes the type and flexibility of the matrix,

the type and configuration of the reinforcement, and the rei nforcement

volume fraction. At the very least, it should be possible to establish

meaningful lower bounds on such properties. Special properties, such as

radiation resistance, could continue to be associated with controlledspecification products that could be incorporated into the basic coding

system

.

It is beyond the scope of this paper to suggest a definitive

system, but it is worthwhile to consider the general form that such a

coding or classification system might take. One possibility, paralleling metals practice, is illustrated in Fig. 3. Here, the first

484
letter(s) and number define the matrix and reinforcement class, identifying the composite as being in a general category, such as epoxyglass or polyimide-graphi te. A following letter defines the reinforcement configuration, e.g., Ell) would define a uniaxially reinforced,

epoxy-glass composite. The next two digits define the nominal fiber

volume fraction. A ± 3% range would be sufficiently precise and would

be compatible with existing manufacturing practice. An El 1)60 coding

thus defines a 60 ± 3 % volume fraction uniaxial glass-epoxy laminate.

The performance of such a laminate could be quite definitively established by adding general information on the degree of matrix flexibility

and specific information on the type of reinforcement. Thus, a designation El U60U-E would not only define a 60% volume fraction uniaxial Eglass reinforced, fully reacted epoxy-matrix laminate but would simultaneously establish meaningful intrinsic performance bounds on the

laminate performance. Extrinsic factors related to laminate quality,

such as void content and interply bond integrity, will affect performance, but such factors must be independently assessed for all laminates.

The fabric used to reinforce laminates is frequently unbalanced,

usually having a lower fiber content in the fill direction of the weave

than in the warp direction. For example, a type of 7628 glass weave

commonly used in the United States has a fill /warp ratio of 32/42. In

the scheme of Fig. 3, information on this unbalance is provided by a

two-digit number defining the fill/warp fraction, or 76 in the cited

case.

This type of system offers flexibility in establishing coding at

the specific level required for the application of interest. The cryogenic industry would probably be reasonably well served by the categories listed in Table I. The list could be expanded or individual

485
designations could be made more or less specific as need arises. Some

cryogenic data already exist for most of the materials on this list,

although in no case are the data sufficient for establishment of statistically significant lower performance bounds.

A standardization scheme of this type would simplify materials

selection and purchase; furthermore, it would minimize costs by providing the maximum flexibility to the laminating industry in meeting the

requirements

.

Consider, for example, the purchase of a G-10 type of laminate

widely used in superconducting magnet construction in the United States.

Under the present system, the cryogenic performance of this type of

laminate purchased on the open market can vary by 30%, if produced by a

reliable laminator, or more, if produced by marginal fabricators.

Specifying the CR grade would solve the problem for critical applications; however, the premium paid for this grade may not be justified for

many applications. Here, the illustrated classification system provides

a solution by allowing the purchaser to specify an E1F48U-E/76 laminate

if the generic equivalent to the CR grade was desired or to specify an

El F40U-E/7 6 laminate as a minimum requirement if greater property variability is acceptable. An E1F40U-E/XX specification would result in

still further economy by removing the requirement for a specific fill/

warp ratio.

SUMMARY

The author has described the ongoing efforts in the United States

to provide a commercial supply of specific control 1 ed-performance laminates required to meet the short-term needs of the cryogenic industry.

486
This approach also provides the basis for a systematic materials development program required to meet future needs, but the current effort

addresses only the most demanding usage of nonmetallic laminates in

cryogenic construction. The overall long-range requirements of the

cryogenic industry justify consideration of a more comprehensive system

of nonmetallic laminate classification that would simplify materials selection while minimizing costs. An example of such a system is provided

to illustrate its potential benefits to both the consumer and the producer .

ACKNOWLEDGMENT

This work was supported by the Office of Magnetic Fusion Energy,

United States Department of Energy.

REFERENCES

1. D. Evans, in Advances in Cryogenic Engineering , Vol . 26, Plenum

Press, New York (1980), p. 83 (in publication).

2. M. B. Kasen, G. R. MacDonald, D. H. Beekman, and R. E. Schramm, in

Advances in Cryogenic Engineering , Vol. 26, Plenum Press, New York

(1980), p. 235 (in publication).

3. R. H. Kernohan, R. R. Coltman, and C. J. Long, Radiation Effects

on Insulators for Superconducting Magnets , 0RNL/TM-61 93 , Oak Ridge

National Laboratory, Oak Ridge, TN (1978).

4. R. H. Kernohan, R. R. Coltman, and C. J. Long, Radiation Effects

on Organic Insulators for Superconducting Magnets , 0RNL/TM-7077

,

Oak Ridge National Laboratory, Oak Ridge, TN (1979)

.

5. B. S. Brown, Radiation Effects in Superconducting Fusion Magnet

Material s

,

submitted to J. Nucl . Mater. (1 980).

487
Table 1. Laminate Classifications of Primary Interest to

Cryogenic Technology

Laminate Category Suggested Coding

Uniaxial Glass-Epoxy E1U55U-E*

El U55U-S2

E1U55U-S

Glass Fabric-Epoxy El F48U-E/76*

El F40U-E/76

El F40U-E/XX

Glass Mat-Epoxy E1M35U-E*

Glass Fabric-Polyester PEI F45U-E/76

PEI F40U-E/76

PEI F40U-E/XX

Glass Mat-Polyester PE1M35U-E

Glass-Polyester Pultrusion PEI F45U-E/1 1

( P )*

Cotton Fabric-Phenolic PH5XXU

Glass Fabric-Melamine Ml F45U-E/76

Ml F40U-E/XX

Uniaxial Aramid-Epoxy E2U60U-49

Aram id Fa brie- Epoxy E2F50U-49/94

Uniaxial, High-Strength

Graphi te-Epoxy

E3U60U-HS

Fabric, High-Strength

Graphi te-Epoxy

E3F60U-HS/1 00

Uniaxial, Medium-Modulus

Graphi te-Epoxy

E3U60U-M

Uniaxial, High-Modulus

Graphite-Epoxy

E3U60U-HM

Uniaxial, 5.6-mil

Boron-Epoxy

E4U50U-56

For illustration purposes only.

Classifications meriting development of special cryogenic grades.

488
List of Figures

Fig. 1 . Flow chart depicting a proposed approach to meeting immediate

needs of nonmetallic laminate development for the MFE , MHD,

and rotating cryogenic machinery industries. Dashed lines

depict the present state of development of the G-l OCR and

G-11CR grades of industrial laminates.

Fig. 2. Suggested flow chart for development of radiation-resistant

nonmetallic laminate insulators.

Fig. 3. Possible coding system for classifying nonmetallic composite

laminates

.

Tabl e 1 .

List of Tables

Laminate Classifications of Primary Interest to Cryogenic

Technol ogy

.

489
Tentative

Specification co

CO

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491
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r

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NONMETALLIC MATERIALS:

RESEARCH AND CHARACTERIZATION

493

THERMAL CONDUCTIVITY OF G-IOCR AND G-11CR INDUSTRIAL

LAMINATES AT LOW TEMPERATURES (PART III)

J. G. Hust and R. Boscardin

Thermophysical Properties Division

Continuous Process Technology Programs

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Thermal conductivity measurements from 2 to 300 K were made on the resins

used in two cryogenic grades of fiberglass/epoxy composites, G-IOCR and G-11CR.

Effects of specimen thickness and inhomogeneity of the bulk sample were

studied.

Also, the thermal conductivity of a European glass-fiber/epoxy composite

was measured to assess the variation between measurements at CERN and at NBS.

The accuracy of the NBS measurements was checked by measurements on a standard

reference material.

495

INTRODUCTION

Extensive low- temperature electrical, thermal, and mechanical properties

characterization of cryogenic grades of G-10 and G- 1 1 fiberglass epoxy composites are being performed at the National Bureau of Standards. These special

laminate grades, referred to as G-10CR and G-11CR, meet NEMA Grades G-10 and

G- 11, MIL-P-18177C, and federal specification LP-5Q9.

In a previous report (Part I),* thermal conductivity data were presented

for G-10CR and G-11CR at temperatures from 2 to 300 K for heat flow in both the

normal and warp directions. These data indicated that the two laminates, produced by a single manufacturer (manufacturer A), had the same conductivity

within about 10%. Also, it was found that the warp direction conductivity was

about 20% higher than that in the normal direction.

In Part II of these reports,^ data were presented showing the differences in thermal conductivity on specimens obtained from another manufacturer

(B) and the effect of large strain. Preliminary data were also presented for

thermal conductivity of the resins used in G-10CR and G-11CR. Further measurements on these resins and the conclusions are presented here.

In addition, this report (Part III) contains data on a European glass

fiber/epoxy composite (Vetronite* 432 .81 )

. ^ The purpose of these measurements is to clarify the cause of apparently large differences between NBS and

CERN data on materials expected to be quite similar. The Vetronite was measured at both laboratories for direct comparison on a single sheet of material.

The NBS measurements were also performed on various thicknesses to illustrate

the absence of significant systematic errors due to interfaces and radial heat

losses.

*Tradenames are used to describe adequately the materials studied. No endorsement of a particular product by NBS is implied.

497
MATERIAL AND SPECIMEN CHARACTERIZATION

G-IOCR resin is a heat-activated amine- catalyzed bisphenol A, solid-type,

epoxy resin. G- 1 ICR resin is an aromatic amine bisphenol A liquid-type epoxy

resin. The G-IOCR resin has an uncured density of 1.10 + 0.01 g/cm3 and the

G- 1 ICR resin has an uncured density of 1.04 + 0.01 g/cm3

.

The resin specimens were cut from blocks supplied by the manufacturer.

The dimensions, masses, and calculated densities of these specimens are given

in Table 1. The resin density calculated from data for the composites (reported earlier) is 1.25 + 0.01 g/cm3 for both resins. Note that the measured

densities of the G-10CR and G- 1 ICR resins are 1.180 g/cm3 and 1.227 g/cm3

,

respectively. The specimens designated as G-10CR resin (a) and (b) are the two

obtained by cutting the original specimen in two, to study the effect of thickness and inhomogeneity.

A sheet of the Vetronite material, measured by K. D. Peterson of CERN, was

obtained and twelve square specimens were cut from it. These were fabricated

by cutting two 1.9 cm wide strips from the supplied sheet. Each of the two

strips were further sectioned into six equal pieces. The twelve resulting

pieces were machined to their final dimensions. Their dimensions, masses, and

densities are given in Table 1. Measurements were performed on selected groups

of these specimens to obtain thermal conductivity values as a function of

thickness.

EXPERIMENTAL PROCEDURE AND DATA ANALYSIS

The thermal conductivity of the test specimens is determined using a thermal conductivity integral apparatus. This apparatus is described in detail by

Hust and Arvidson. 4 The estimated uncertainty of the results is 5%, unless

498
otherwise noted, as verified by measurements on NBS standard reference material

SRM 735.

The data range is normally 2 to 300 K and the data are normally spaced

approximately logarithmically within this range. The measured thermal conductivity integrals are fit by a special least squares technique (see Hust and

Arvidson [1]) to yield the parameters for the following equation:

n

i K (T) = E A.[ln(T + 1)] Eq. 1 i=l

1

RESULTS AND DISCUSSION

To facilitate future analysis of composite data, measurements were performed on specimens fabricated from the pure resins used in G-10CR and G-11CR

(see Table 1). These specimens were prepared from cured pieces supplied by

the manufacturer. Twenty-one data points between 2 and 304 K were obtained on

the original G-10CR resin specimen and sixteen data points were obtained from 4

to 304 K on the G-11CR resin specimen. The results, essentially the same within the precision of the data (+4%) for both resins, are illustrated in Figure

1. These values, however, are more uncertain and imprecise than usual because

the specimen thicknesses were too large for optimum accuracy. Because of this

potential problem with these first sets of data and because of the possibility

of thermal radiation effects in these nearly transparent specimens, it was

decided to section the G-10CR resin specimen to produce two thinner specimens.

The thinnest specimen, about 1/3 the original thickness (see Table 1), was

measured at 18 temperatures from 4 to 300 K. The results, appreciably lower

(25%) than the original G-10CR resin data, are illustrated in Figure 1. This

large difference was not understood and led to a series of measurements near

499
room- tempera ture . The results of the room temperature measurements are given

in Table 2.

The conclusion reached from these measurements is that there is a significant inhomogeneity in the G-10CR resin specimen. The source of this inhomogeneity is uncertain but it is likely to be found in the original resin-mixing and

block-curing procedure. The results of the measurements on specimens #1, (a),

(b) and (a + b) suggest that the thermal conductivity is greater near the

middle of the specimen, since both the results of the original measurement on

the whole specimen and the later measurement on the reassembled specimen (a +

b) were higher than (a) or (b) measured separately. It is also noted that the

reassembled specimen resulted in an intermediate value, which is consistent

with the above conclusion, since the center section was partly removed by

cutting. The average thermal conductivity of nonhomogeneous specimens can be

direction dependent, i.e., the results can depend on whether the hot face is on

the low-conductivity side or the high-conductivity side. This effect is not

quantified, because the orientation of the specimen in the chamber was not

recorded in all cases. Whether the G-11CR resin specimen contains similar

inhomogeneities is unknown because it was not sectioned. Since the percentage

difference in thermal conductivity between the small and large G-10CR resin

specimens is essentially independent of temperature, the thermal radiation

component to total heat transfer is small compared with the conduction

component.

One Vetronite specimen was measured near room temperature for comparison

with the CERN data. In addition, as a check on the possibility of systematic

errors due to interface resistance, radial heat losses, and the calibration

constant of the probe, measurements were performed on randomly selected pieces

500
in the form of stacks (2 thick, 4 thick, and 12 thick). The results of the

measurements are given in Table 3. These data reveal only a very small

systematic deviation as a function of thickness. The constancy of the results

over the extremes of thickness used here illustrate that boundary layer

resistance (at the smallest thickness) and radial heat losses (at the largest

thickness) are not significant experimental errors. In all cases, these

results are about 45% higher than the results from CERN. This confirms that

the differences are truly experimental in nature and not specimen differences.

The possibility of systematic errors in the NBS thermocouple calibrations and

other instrumentation calibrations has been ruled out by measurements on SRM

735. 4

501
REFERENCES

1. J. G. Hust, "Thermal Conductivity of G-10CR and G-11CR Industrial Laminates at Low Temperatures," in Materials Studies for Magnetic Fusion

Energy Applications at Low Temperatures - H_, NBSIR 79-1609, F. R. Fickett

and R. P. Reed, Eds., National Bureau of Standards, Boulder, CO (1979),

pp. 417-436.

2. J. G. Hust, "Thermal Conductivity of G-10CR and G-11CR Industrial Laminates at Low Temperatures (Part II)," in Materials Studies for Magnetic

Fusion Energy Appl i cations at Low Temperatures - III , NBSIR 80-1027, R. P.

Reed, Ed., National Bureau of Standards, Boulder, CO (1980), pp. 395-404.

3.

K. Dahlerup-Petersen and A. Perrot, "Properties of Organic Composite

Materials at Cryogenic Temperatures: CERN ISR-BOM/79-39." (November

1979).

4.

J. G. Hust and J. Arvidson, "Thermal Conductivity of Glass Fiber/EPOXY

Composite Support Bands for Cryogenic Dewars," unpublished report

275.03-78-2, National Bureau of Standards, Boulder, CO.

502
TABLE 1

DIMENSIONS, VOLUMES, WEIGHTS, AND DENSITIES

OF THE THERMAL CONDUCTIVITY SPECIMENS

DIMENSIONS

(cm)

SPECIMEN x

a x

b

1 2

x

a

3

VOLUME

3

(cm )

WEIGHT

(g)

DENSITY

3

(g/cm )

G-10CR RESIN #1 1.915 1.913 1.781 6.524 7.698 1.180

(a) 1.915 0.652 1.781 2.224 2.615 1.176

(b) 1.915 1.041 1.781 3.550 4.175 1.176

(a+b) 1.915 1.693 1.781 5.774 6.790 1.176

G-11CR RESIN 1.915 1.915 1.915 7.023 8.620 1.227

VETRONITE 432.81 (Al) 1.786 0.199 1.770 0.629 1.200 1.91

(single pieces) (A2) 1.780 0.198 1.768 0.623 1.194 1.92

(A3) 1.783 0.196 1.768 0.618 1.189 1.92

( A4) 1.780 0.196 1.765 0.616 1.186 1.92

(A5) 1.780 0.198 1.765 0.622 1,190 1.91

(

A

6 ) 1.780 0.199 1.765 0.625 1.193 1.91

(Bl) 1.780 0.201 1.768 0.632 1.202 1.90

(

B

2 ) 1.775 0.200 1.768 0.628 1.200 1.91

( B3) 1.778 0.197 1.765 0.618 1.189 1.92

( B4) 1.775 0.196 1.763 0.613 1.183 1.93

( B5) 1.770 0.198 1.760 0.617 1.180 1.91

(

B

6 ) 1.768 0.200 1.753 0.620 1.176 1.90

(stacked groups)

A3 & B3 1.780 0.393 1.766 1.235 2.328 1.92

A3, A5, Bl, and B3 1.780 0.792 1.766 2.490 4.770 1.91

all 12 specimens 1.778 2.378 1.765 7.462 14.282 1.91

a - Values of and X 3 for the stacked groups are averages of the

dimensions of the individual pieces.

b - Values of X

2 for the stacked groups are the sums of the thicknesses of

the individual pieces.

503
TABLE 2

MEASURED THERMAL CONDUCTIVITY VALUES AT 280 K OF THE G-10CR

RESIN SPECIMENS [SPECIMEN #1 REFERS TO THE ORIGINAL SPECIMEN, (a)

AND (b) ARE THE SPECIMENS OBTAINED AFTER SECTIONING SPECIMEN #1]

Specimen Identification and Thickness (cm)

Thermal Conductivity

(W m"

1 K -1

)

G-10CR Resin #1 1.913 0.296

G-10CR Resin (a) 0.652 0.230 + 0.008a

G-10CR Annealed 15 h 0 100°C (a) 0.630 0.232

G-10CR Resin (b) 1.041 0.246

G-10CR Annealed 15 h 0 100°C (b) 1.026 0.238

After Anneal i ng (a + b) 1.656 0.252

G- 1 ICR Resin 1.915 0.282 + 0.005^

G- 1 ICR Annealed 15 h 0 100°C 1.915 0.278

a Range of 4 measured

b Range of 2 measured

val ues

val ues

504
TABLE 3

THERMAL CONDUCTIVITY VALUES FOR VETRONITE 432.81 NORMAL

TO THE FIBERS AS A FUNCTION OF THICKNESS AT 284 K.

Specimen Identification

(see Table 1)

Thermal Conductivity

W m"

1 K"1

B3 0.447

B3, A3 0.450

B3, A3, Bl, A5 0.451

All 12 Specimens 0.480
300 CM ro

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FRACTURE ANALYSIS OF COMPRESSION DAMAGE

IN GLASS FABRIC/EPOXY, G-IOCR, 295 K to 4 K

Betty A. Beck

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

ABSTRACT

This report describes scanning electron and optical microscope fractographic and modeling studies of G-IOCR compression fracture at room and cryogenic temperatures. The primary objective of the investigation has been to assess the effect of cryogenic temperature on the failure behavior of these

composite materials. Macromechanical analysis in terms of gross composite

fracture has been designated as primary. Micromechanical (or constituent)

analysis has been specified as secondary. This identification has been used

for descriptive and comparative purposes rather than to indicate the level of

importance in the fracture process. Results of the fractography have shown

that the failure events included three kinds of primary fracture patterns

which, to varying degrees, are temperature-related. These fractures have been

identified as: 1) a 55° slant mode; 2) a chisel or V-notch mode; and, 3) a complex or mixed mode. The influence of constituent properties and behavior

have been integral considerations in assessing the fracture mechanism and in

analytical interpretations of the fracture behavior. As in other materials,

the damage zone in these composites is the result of either discrete or interactive constitutent failure processes (or both) occurring in a highly strained

region of the specimen. Fiber microcracking or fracture, matrix yielding or cracking, and interface failure are examples of such events. Nucleation and

propagation of fracture, as in other fi ber-reinforced composites, is a function

of microstructure, microdeformation, and the structural interaction of the

constituents. Compressive strength data variability has been graphically

correlated with primary and secondary fracture characteristics in order to

understand the low temperature effects better.

507

Contents

Page

Abstract

A. Introduction 5 H

B. Previous Studies 512

1. Introduction 512

2. Constituent Properties and Behavior 515

(a) matrix resin

(b) glass reinforcement

(c) reinforcement/matrix interface

3. Composite Room Temperature Behavior and Failure 518

(a) general

(b) toughness

(c) compression

4. Studies of Glass Reinforced Nonmetal lies at Low Temperature 526

5. In Summary and Concluding Remarks 531

C. Low Temperature Effects on Fracture Modes of G-l OCR Tested in

Compression at 295 K, 76 K, and 4 K. 532

1 . Introduction 532

2. Material Specifications, G-l OCR 534

3. Test Procedures 535

4. Fracture Analysis 535

5. Results and Discussion 540

(a) general

(b) fractographic observations and analysis

(c) conclusions

509
Contents (continued)

Page

6. Recommendations for Future Studies 551

7. References 552

8. Tabl es 557

9. 11 lustrations 561

510
A. Introduction

In general, composite failure is complex because of the interrelated

behavior and properties of the constituents, i.e., the matrix, the reinforcement, the reinforcement/matrix interface as well as the variability resulting

from the fabrication process. Fracture analysis experience has shown that no

single failure mechanism stands out as a primary cause of composite fracture.

Rather, there seems to be a sequence of failure events that lead to total

composite failure (1-7).

Studies have shown that the predominant primary mechanism of composite

fracture as well as the sequence of constituent failure events varies significantly for different composite types (6). Included with this failure pattern

variability are the energy dissipative mechanisms at fracture initiation and

growth which also differ from one material system to another. This lack of

failure behavior continuity and the inherent complexity of fabric-reinforced

nonmetallics emphasize the need for systematical ly building a theory of failure

based in part and initially, on detailed observations of fractured specimens.

Numerous investigators have studied and described failure and fracture

behavior in both discontinuous and uniaxially continuous glass-fiber reinforced

nonmetal lie composite materials (1-7). For the most part analytical interpretations, methods, and models, where they exist, have been developed around

some theoretical or experimental aspect of a composite's room temperature

behavior. Low temperature studies have had a characterization orientation

primarily.

Fabric-reinforced nonmetal lie composites possess a very complex geometry

and behavior, and as such, are more difficult to describe both qualitatively

and anal ytical ly. These materials have, to date, eluded conventional analysis

511
and modeling. Most of the difficulty seems to be in the inherent complexity of

modeling the reinforcement. Because of the dearth of analytical studies for

these materials at both room and low temperatures , there is essentially no

analytical base on which to build a theory of failure.

In this study and through exploratory data analysis, a detailed description of primary and secondary fracture patterns has been developed to elucidate

basic mechanisms of fracture and to determine temperature-rel ated modes of

failure. The results of this investigation will provide a foundation for

developing failure analysis techniques based on correlating fractographic

features with fracture properties and stress analysis at fracture.

B. Previous Studies

1 . Introduction

A summary of information from the published literature has been included

in this report. In order to identify, interpret, and analyze failure patterns

in the G-10CR specimens, the literature review has provided a basis for developing

a fractographic and an analytic frame of reference. Some of the published

results have been modified or extended in part to this composite system which

involved a different reinforcement geometry and type and loading condition.

In utilizing and synthesizing the information on the behavior of matrices,

reinforcements, and to some extent, interfaces, a rationale has been established

for the sytematic analysis of failure in the G-10CR specimens. The state of

the present study would be signficantly less complete without the direction

provided by the earlier studies. Because there are so few published quantitative or even qualitative analyses about failure behavior in fabric reinforced

nonmetal 1 ics , this approach has seemed necessary.

512
Most investigators agree, that, to formulate a better analytical understanding of the fracture process in all composite materials, knowledge, not

only of the predominent mode of failure but of the sequence of failure events

and their probability of occurrence is a primary consideration in the difficult

task of reconstructing the fracture process from initiation through propagation

to final failure (1-3, 6-14). Similarly there is a consensus about the importance of fractographic analysis in advancing a comprehensive theory of failure

and failure analysis techniques (14).

There seems to be a lack of a well-ordered knowledge base that can be

used to develop 1 ow-temperature strength theories of failure as well as predictions for the onset of failure. This has been particularly true for the

glass-fabric-reinforced nonmetal 1 ics. There have been only a few studies of

the 1 ow-temperature behavior of glass-reinforced nonmetal lie composites.

Constituent behavior, the matrix resins, and glass reinforcement have received

somewhat more attention. The structural dependence of Young's modulus, Poisson's

ratio, and strength, for example, has been shown to be rather small in epoxy

resins. On the other hand, their fracture energy shows significant room- and

low-temperature dependence on structure (15). For the composites made up of

these materials, the low-temperature changes in strength, moduli, and other

properties are incompletely understood structurally. Correlation of changes

in static mechanical properties of glass-reinforced-epoxy composites with

macrolevel fracture characteristics at cryogenic temperatures is almost nonexistent.

A prevailing idea among the many studies is the need to relate fracture

modes to structure and states of stress in the development of valid failure

criteria. Morgan et al . , for example, point out that the network structure

and void characteristics of epoxy resins control their mechanical response (16).

513
The network structure, which varies with chemical composition and curing

conditions, imparts rigidity to the resin. Voids act as stress concentrators

and, as such, degrade mechanical response by adversely affecting the effective

shear strength. The structure and properties of the resin are significant to

the load transfer function of the matrix. Low temperatures can adversely

affect this load-transfer mechanism.

Fracture analysis of all composite systems is a function of the deterioration of the unit structure. Glass-reinforced materials are no exception.

The constituents impart a major effect to the breakdown under an applied load.

Rosen, as well as other investigators, has shown that there is no unique

mechanism that will adequately describe failure in all fiber-reinforced nonmetallics (1=3, 4, 5, 9). The individual properties and characteri sties of

the constituents (the matrix, reinforcement, and interface) and their interactions in the process of degradation and failure yield a complex, coupled

behavior in total composite fracture that is difficult to trace to its origins.

The contributions of the variety of interacting mechanisms are not easily

discernible in the complexity of general composite failure (10). To a great

extent, for example, composite strength, which is related to how effectively

the stress can be transmitted from one fiber to the next, is a function of the

individual constituents (17).

Kelly, in discussing uniaxially reinforced composites subjected to an

applied load parallel to the fiber, emphasizes a modulus mismatch that causes

displacements in fiber and matrix to differ and leads to stress concentrations

proportional to the difference in moduli. Shear stresses and strains resulting

from the stress concentrations are the mechanism for load distribution between

matrix and fiber via the interface (17, 18). The major part of the load is

514
carried by the reinforcement if the matrix modulus is such that a matrix

displacement is greater than that of the fiber reinforcement.

Matrix and interface characteristics (void content and shear properties)

are primary to an effective load-transfer mechanism (19). Cracking in the

resin matrix can result in a poorer distribution of the load or abrasion of

the glass. Most investigators agree that the load is transferred by shear

stresses at the fiber/matrix interface and that the rate of load transfer is

limited by the strength of either the interface, the matrix, or the reinforcement, whichever is least (3, 17). A strong interface (low void and crack

content) provides a more effective load-transfer effect than a weak interface

(high void and crack content that degrade bonding and shear strength) (17). The

load-transfer effectiveness of the matrix is adversely affected by increased

resin brittleness, increased Young's modulus, and the existence of residual

stresses in the composite because of resin contraction and shrinkage during

cure (3, 20).

The matrix resin protects the fibers and provides a shear rigidity

between fibers and lamina, strength and stiffness to the reinforcement, and a

coupling effect to the constituents in response to tensile, compressive,

bending, and shear loads. Additionally, it sustains a high level of strain

concentration that occurs as the result of the high-modulus, brittle fibers in

the gl ass-reinforced-epoxy composites.

2. Constituent Properties and Behavior

(a) Matrix Resin

In general, polymer failure occurs by excessive deformation, crazing or

cracking, and tensile or shear fracture. Cross-linked polymers fail by cracking

515
and fracture (21). Structural ly, there is evidence to indicate that stress

concentrations can lead to cracking at the cross-links if the cross-link

length is too short for plastic flow. Shear fracture for polymers can occur

on either a single plane or simultaneously along two planes of maximum resolved

shear stress lying at 45° to the longitudinal axis of testing. Fracture

results when deformation, due to a concentration of large shear strains, is

restricted to the one or two planes of shear failure.

Strength and elasticity of epoxy resins show little or no dependence on

cross-link distance for temperatures in the range above 50 K to less than 20 K.

Energy at fracture, however, for these same resins, is strongly dependent on

cross-link distance, increasing as the distance increases (15).

The response mechanism of a polymer to an applied mechanical stress is a

function of the ability of the molecules to move past or over one another. At

cryogenic temperatures there is an increase in the rigidity of these materials

owing to an increased rigidity in the molecular structure.

Epoxy resins, in general, exhibit room-temperature glasslike textures

typical of brittle materials that have characteristic low-energy-absorbing properties, low viscosity, and low shrinkage when cured (5, 22, 23). They don't

form bubbles, which minimizes buildup of internal stresses in the cured resin.

(b) Glass Reinforcement

The high-strength, stiff glass fibers, in the form of a rigid molecular

network with no mechanism for ductility, are subject to a statistical distribution of flaws and defects characteri stic of brittle materials. As such,

fiber failure is by random microcracking and fracture inherent to brittle

516
failure. An accumulation of fiber breaks is considered by many investigators

to be a major mechanism in total composite failure. A strong, stiff matrix

contributes to a greater probability of a statistical failure mode of the

glass fibers (1-3, 7, 8, 11).

(c) Reinforcement/Matrix Interface

Debonding of the fiber from the matrix and delamination of one lamina

from another are interface failure modes. Most composite studies have shown

that the fiber/matrix interface is primary to load distribution from matrix to

fibers. Strength, mode of failure, and work at fracture are a function of the

behavior of the interface. Interface shear strength is a function of the

density of voids formed during fabrication. High void content results in low

shear strength and leads to a weak interface that can be split by stresses set

up ahead of an advancing crack (17).

There is some evidence to indicate that fiber/matrix debonding and matrix

cracking can interact to reduce load concentrations and result in an increased

composite stress. Studies have also indicated that, where there is an increase

in composite strain at fracture, it may be due to deterioration of the fiber/matrix

i nterface.

In synthesizing fracture information for glass-reinforced-plastic composite systems, constituent characteristics and their highly coupled nature

and behavior relative to the failure process must be integrated with the

analysis. The degradation and failure mechanisms may occur independently,

simultaneously, or consecutively in total failure.

Energy-dissipative mechanisms and levels thereof are integrally important

in sequence analyses of failure events (4, 24, 25). Crack branching, for

517
example, in fiber-reinforced-composites extends the size of the fracture

surface and results in a large dissipation of energy. Debonding, for brittle

fibers enclosed by a semibrittle matrix resin, seems to occur first along an

interface, and the debonding work done is stored in the fiber as elastic

strain energy. When a debonding sequence is completed, the fiber pulls out of

the matrix. This occurs because the load on the fiber, in order to split it

from the matrix, is higher than the load for pullout. As a result, the energy

dissipated at fiber pullout is higher than the work done in debonding (4).

3 . Composite Room Temperature Behavior and Failure

(a) General

Primary, considerations in composite fracture and fractographic analyses

include the identification and description of: (1) the predominant failure

mode of the composite; (2) the failure modes of the constituents; (3) their

probability of occurrence in the failure of a specific composite material

under a well-defined loading condition; and, (4) a probable sequence of failure

events from initiation to propagation and final fracture. Collins et al .

,

for

example, conclude that, in determining which mode dominates failure behavior,

fractographic analysis is a function of which model is assumed (11). Their

study identifies and summarizes two approaches: (1) a cumulative fracture

propagation model and (2) a noncumulative model. These models have also been

described and expanded upon in other investigations (2, 9).

The cumulative model assumes that the total applied load is borne by the

fibers. The matrix acts only as a mechanism for redistribution of the stresses

when a fiber breaks locally. This model is based on a statistical strength

distribution due to inherent defects characteristic of the brittle glass

513
fibers. Total composite fracture occurs after enough fiber fractures

have resulted in a weakened cross section so that it can no longer

support the applied load. This model is valid if an assumption is made

either that (1) the load is distributed uniformly over the remaining

fibers after one breaks or (2) surviving fibers are subjected to static

and dynamic load concentrations. The noncumulative mode identifies

composite failure as a function of crack propagation through the matrix,

interface, and reinforcement. As a result, the failure loads are considerably

below those which would cause a significant number of fibers to break.

According to these investigators, fractographic analysis will be

approached differently depending on which mode is assumed. The cumulative

mode implies, for example, that interface debonding of fiber and matrix

and fiber pullout result in an ineffective fiber length, which carries

little of a tensile load in a failed area. The noncumulative mode, on

the other hand, implies that interface debonding and fiber pullout

determine an optimum path of crack propagation through the matrix,

including crack branching at the interface.

Zweben has determined that such factors as fiber/matrix debonding,

interface cracking, matrix plasticity and cracking, and the three-dimensional

nature of the material act as energy-absorbing mechanisms that tend to

reduce load concentrations, effectively increasing composite stress. On

the other hand, a lack of uniformity in the spacing as well as cracks in

the fibers tends to increase stress concentrations and decrease composite

failure stress (2, 9, 11, 24).

Early studies of the tensile failure of a uniaxially reinforced

matrix identified a simple failure model based on a uniform strain

throughout the composite. As a result, fracture occurred at the failure

strain of the fibers alone (2,7). More recently most investigators

51 9
have postulated a nonuniform strain distribution, which has directed

consideration of the influence of flaws in the brittle reinforcement on

composite failure (4, 5, 7, 9).

Rosen, and others, have developed a rather comprehensive description

of failure in fiber-reinforced composites (2, 6, 7). This model assumes

that because of flaws inherent to the brittle fibers, initial fracture

will most likely occur in the fibers causing a local perturbation of the

stress field. The way the load is redistributed near the broken fiber

ends and the resulting stress concentrations may result in: (1) separation

along an interface (i.e., fiber debonding or lamina-to-lamina delamination);

(2) fracture of nearby fibers (since they are subjected to a greater

stress intensity than those at some distance); (3) propagation of transverse

cracks; (4) immediate composite fracture; or (5) no further fracture.

Because the strengths of the brittle fibers are not uniform, they will

break randomly throughout the material as the applied stress is increased.

The fractured fibers are then reloaded to some level before being pulled

out of the matrix. There is, however, a critical length at which the

fibers will no longer support the load. As fractures accumulate, the

end result is a number of fibers of ineffective lengths interacting to

produce a weakened cross section of the composite, which leads to total

failure. Since the shortened fibers can no longer effectively transmit

the increasing load, composite fracture occurs as the maximum shear

strength of the matrix is exceeded. A strong, stiff matrix, the nonuniform

strength of the fibers, and a statistical accumulation of flaw-related

fiber fractures are the controlling factors in the failure characterization

of this model. Because the fibers do not have uniform strengths, strong

fibers will carry an increasingly greater load as weaker fibers break.

Additionally, the average strength of some of the individual fibers may

520
be greater than the coupled strength of fiber bundles, which can significantly affect the fracture sequence. Because of the lack of uniformity

in the glass-fiber strengths, investigators have utilized the Wei bu 11

distribution as statistically descriptive of the strength variability

(4, 5, 7, 9, 26).

There is a distinction between the failure processes (including

stress distribution and level) that control the onset or initiation of

failure and those that control the fast propagation of a crack. The

former is more difficult to reconstruct fractographically and has thus

received less study. Because of the variety of micromechanisms of

failure occurring and interacting, failure mechanisms at the onset of

damage must be distinguished from those close to fracture in terms of

stress levels and states. Rosen feels that, in developing a failure

criterion, it's initially important to understand internal irregularities

in the state of stress (7).

Tetelman has determined that composite fracture initiates either

by: (1) fiber fracture, (2) interface shear fracture (debonding or

delamination), or (3) matrix tensile or shear fracture, depending on

which of the above occurs first at a lower value of applied strain.

Final fracture may be the result of some or all of the above mechanisms

(3).

Mull in has analyzed energy-absorbing mechanisms in high-modulus fiber

composites (24). His analysis indicates that, at a fiber fracture site,

strain energy will be increased in adjacent fibers by a stress-transfer mechanism acting through the interface and the matrix. The localized redistribution of load can produce either matrix cracking perpendicular to the fiber

or debonding at the interface. The probability of the former mechanism occurring

reflects a strong interface bond; the latter, a weak one. Debonding can

521
result in a localized sliding of the fiber in the interface socket accompanied

by fiber fracture and pullout. Beaumont has identified postdebond sliding,

rather than fiber fracture and pullout, as a major toughening mechanism in

fi ber-rei nforced composites (27). Vinson and Chow associate fiber pullout

with the flaws in the brittle fibers that lead to nonuniform strengths (4).

They assume that the matrix has cracked completely. As a result, fibers in

the cracked regions become stressed more than the remaining ones, which are

firmly fixed in the matrix.

Zweben cites a study of composite failure that associates fracture

sequence with a propagating stress wave that fractures the fibers destroying

their 1 oad-carrying capability (25). The matrix, as a result, cannot sustain

the applied load level without the strengthening effects of the reinforcement.

This leads to total composite failure with most of the energy being dissipated

by the matrix.

Additional models for describing failure initiation and propagation with

associated energy dissipative mechanisms have been proposed. One of these

indicates initiation is the result of debonding brought about by increasing

fiber fracture strength or decreasing matrix or interface shear strength (3).

Voids introduced during fabrication degrade interfacial shear strength and

significantly influence debonding. Deformations that develop during initial

impairment of a fi ber-reinforced material can lead to a tensile fracture of

the reinforcement, which is most probably due to bending followed by a final

shear mode of composite failure (5, 27). Materials with a high void content

can fail by crack extension parallel to the fibers, whereas those with low

void content and high fracture toughness have a characteristic shear final

failure mode. Residual stresses and strains introduced during fabrication are
the result of differential thermal expansion characteristics between reinforcement and matrix. The existence of these stresses can complicate the analysis

of failure initiation. If, as some investigations have assumed, failure first

occurs in the brittle matrix and the crack propagates toward the fibers, it

can propagate across the fibers if the interface is strongly bonded and continue to brittle failure. If, however, the propagating crack approaches a

weak interface, debonding of fiber/matrix interface occurs. Debonded fibers

then fracture as a function of the defect distribution in the imperfect fibers.

Stress concentrations ahead of an advancing crack can cause the interface to

fail. When a propagating crack reaches a fiber, it is diverted along the

cracked interface and blunted. In order for the crack to continue to propagate, it must be reinitiated by increasing the load. Some investigators have

emphasized that this sequence of events is the primary toughening mechanism in

fiber-reinforced nonmetal lies (4, 5, 27).

(b) Toughness

Toughness in a metal is the ability to absorb energy during plastic

deformation without fracturing. In a static tensile test this energy can be

measured by the area under a stress-strain curve, which represents the work

required to fracture a specimen.

In glass-reinforced epoxy composites there is no direct toughness analogy

with a metallic system because of the plastic nature of the metal. If, however, toughness is considered to be the ability to redistribute stresses in

the presence of a notch to reduce the effects of high stress concentrations by

dissipating the energy, then the work required to fracture a composite specimen

should represent toughness properties analogous to a metal. Toughness in a

523
glass-fiber-reinforced plastic seems to be a function of the interfaces and

their ability to deflect cracks propagating normal to them (4, 5, 12, 17, 22,

27). If there is a strong interfacial bond, there is a high probability that

stress concentrations set up by failure of one fiber will nucleate cracks in

nearby fibers leading to progressive fiber failure. If the bond is weak,

interfacial shear strength will be degraded and stress concentrations ahead of

a crack advancing normal to the fibers can split the interface. As the crack

reaches a flawed boundary, it will be blunted or diverted parallel to the

interface. For the crack to continue to propagate, it must be reinitiated by

a sufficient stress buildup as the applied load is increased. If this process

can be repeated as the cracking spreads, large amounts of energy can be absorbed.

Some investigators have indicated that toughness relates to the energy

dissipated as fiber/matrix debonding occurs. The toughness mechanism becomes

a sliding action between fiber and socket, accompanied by friction effects and

eventual fiber pullout (24, 27).

Fiber-reinforced nonmetal lies are considered to be tough and not notch

sensitive.

(c) Compression

Compressive stresses are the result of either an externally applied load

or are due to shrinkage stresses developed during cooling because of differential thermal expansion characteristics between matrix and fiber. Rosen and

others have developed models for uniaxial continuous-filament composites

subjected to an externally applied compression load parallel to the fibers (1,

7, 28-30). The model postulates that failure occurs by microbuckl ing of the

fibers on an elastic foundation. The analysis identifies two fiber buckling

C24
modes, a shear mode and extension mode, each descriptive of matrix

deformation. For this model the matrix shear modulus controls composite

compression strength. Buckling can be prevented by a stiff matrix and

strong bonding at the interface (5, 7, 20, 31, 32). There seems to be

little evidence to show whether, under compressive loading, local buckling

precedes and causes crack propagation or follows crack initiation and

propagation (5).

There is a study cited in Broutman that investigated a linear

relation between interlaminar shear strength and compressive strength

for flat filament-wound panels. That analysis showed compression strength

as a function of resin properties and either high interlaminar or horizontal

shear strengths (5)

.

For his investigation on compressive fracture Chaplin has expanded

the development of the shear mode to a shear instability in a volume of

material rather than failure in a plane as a function of a resolved

shear stress. The shear instability produces significant shear deformation

in the matrix and bending of the reinforcement (32).

Externally applied compressive loads parallel to the fibers in a

unidirectional fiber-reinforced composite can result in failure by:

fiber microbuckl ing in a shear mode; tensile or compressive fiber fracture;

shear failure of the matrix; interfacial failure due to delamination or

debonding or both; or compression-induced bending accompanied by secondary

tensile and shear stresses (4, 5, 7, 31-33). Fiber pullout and fracture

are failure events in the debonding sequence. A major characteristic of

composite shear failure is lack of damage to the fibers in the direction

of shear.

A compressive stress will not propagate internal cracks nor will it

produce crazing, a failure phenomenon characteristic of many matrix resins.

Rather, a compressive stress normal to a crack tends to close the crack.

525
Secondary tensile and shear stresses, however, will propagate internal cracks

and are significant to the resulting fracture behavior.

Ryder and Black, in discussing types of compression testing, indicate

that there are several possible compression failure modes in fi ber-reinforced

composites. Each has a different ultimate strength. For each testing procedure, failure occurs at the lowest strength with a corresponding failure

mode. The range of possibilities suggests that compressive strength for these

materials is not a precise term, but rather, is one of definition (34).

Extensive work on static compressive behavior of glass-epoxy composites

has shown that test results and the failure mode are sensitive to the method

of gripping (misalignment causing the specimen to bend), type of reinforcement,

and fiber volume fraction. Three major compressive failure modes have been

identified: (1) fiber buckling, (2) shear, and (3) interfacial, with the

shear mode predominating (7, 28, 33). Total brittle compression fracture of a

composite specimen results in separation of the specimen into two or more

parts with little or no measurable inelastic deformation.

If deformations develop initially, they can lead to either (1) tensile

fracture of the reinforcement due to bending with final fracture of the

composite in a shear mode or (2) shear displacement in the matrix, which can

cause stress concentrations and crack extension along the fibers if voids are

present in a contiguous area.

4. Studies of Glass Reinforced Nonmetallics at Low Temperature

Fiber and matrix characteristics have a significant influence upon the

mechanics of deformation and fracture of fiber reinforced nonmetallics at both

room and cryogenic temperatures . As such, temperature-rel ated behavior and

526
properties of the constituents are important in the analysis of temperature

effects on failure modes for these composites.

Polymer behavior is a function of a molecular chain response mechanism,

which can experience a temperature-related change as temperatures are decreased

to those representing cryogenic environments. Such change can have strong

effects on material properties and fracture behavior. At room temperature the

molecular chain is in thermal motion around an equilibrium position. At this

temperature the molecular response to an external stress (that is, the coiling

and uncoiling of the chains and their moving over and past one another) is

easier because there is no fixed molecular position. In this configuration,

the polymer chain provides an energy-absorbing mechanism for stress waves

transmitted through the material. As the temperature is decreased, the thermal

energy of the molecules decreases and the elastic mobility of the polymer

chain slows down. As a result, the chain becomes more rigid, the energyabsorbing alternatives become more restricted, and attractive forces become

more effective in the resistance of materials to deformation. Under an applied

load the decreasing ability to dissipate energy through the molecular structure

leads to an increased resin rigidity until the cohesive strength of the

polymer has been exceeded, ultimately leading to fracture as a dissipative

mechanism. This molecular rigidity results in a stronger, stiffer material at

low temperatures because the polymer chains are unable to slide over one

another effectively (3, 15, 16, 23. 35).

For epoxy resins, there are cross-linking effects in the molecular

structure as well as the effects of molecular anistropy of the binding forces

owing to the directional nature of the primary covalent bonds and secondary

527
Van der Waals forces, which act along and across the molecular chains. The

cross-linked structure imparts some degree of rigidity to the polymer.

Epoxy resins do not seem to be subject to changes in structure with

decreasing temperature. They exhibit elastic-plastic stress-strain behavior

at room temperature but have nearly linear stress-strain characteristics at

low temperature. Strength (tensile, compressive, and shear) and modulus

increase with decreasing temperature. Fracture strain, a function of molecular cross-linking, approximates 2% at 77 K and 4 K.

Softer and Mol ho tested a Bisphenol-A epoxy resin at cryogenic temperatures (77 K and 20 K) . Their results indicate that, in comparison with roomtemperature properties, (1) there is an increase in strength and modulus in

tension and compression; (2) there is a change in elastic-plastic behavior at

room temperature to essentially linear elastic behavior at 77 K and 20 K; and

(3) there is a decrease in fracture strain and resin toughness (as measured by

the area under the stress-strain curve) with decreasing temperatures. The

rate of decrease for (2) and (3) moderates between 77 K and 4 K. The resin

exhibited a maximum notch toughness (the capacity of the resin system to sustain loads in spite of defects in the matrix) value at 77 K and decreased to

below the room temperature value at 20 K. Their assessment of the fracturesurface examination of this particular resin specimen indicated a smooth,

almost glasslike texture typical of low-energy, absorbing, brittle materials (23).

Hartwig's low- temperature analysis indicates, in general, that the chemical structure of epoxy resin chains, including cross-link density, does not

have a marked effect on elastic mechanical parameters. Poisson's ratio seems

to be entirely independent of structure. Energy at fracture, on the other

hand, decreases with higher cross-link density at cryogenic as well as higher
temperatures (15). Other investigators have noted a relation between the

degree of cross-linking and void density and their detrimental effects on

matrix and interface shear strength. This effect, however, was not investigated below room temperature, probably because 1 ow-temperature factors were

not a part of the study (21).

Kadotani determined in his low- temperature studies that the lower the

glass transition temperature, T^, of a resin, the larger its tensile elongation at 77 K. The larger the elongation at 77 K, the higher the strength and

elongation of the composite at 77 K (36).

Glass materials used for reinforcement are brittle at room temperature

and below. Their rigid silica network structure and lack of molecular order

is such that there is no inherent structural mechanism for ductile behavior

characteristics. Because of their brittle nature, strengths are not uniform

and their failure modes at room and cryogenic temperatures reflect these

brittle characteristics.

Glass-fabric-reinforced epoxy composites exhibit significant increases in

tensile, compressive and shear strength, tensile modulus, Poisson's ratio, and

tensile strain to failure with decreasing temperature compared with roomtemperature values (20, 36, 37). There is considerable evidence to indicate

that

:

1. A stiffer resin holds the reinforcement more firmly fixed for loading

parallel to the principal direction of fabric-reinforced composite.

As a result, there is a greater resistance of the fibers to displacement or deformation.

2. Fi ber-to-fi ber progressive fracture leading to final composite

fracture provides an enhanced energy-absorbing mechanism, the

529
characteristics of which reflect the sum of the energies of the

individual fibers. As a result, more energy can be absorbed at room

and low temperatures.

3. The increase in strength at low temperatures is the result of interface failure and crack blunting, a progressive mechanism that

imparts toughness to the material.

4. The high strain of the composite at fracture at the lower temperatures is a direct function of the low-temperature, high-fracture

strain characteristics of the glass reinforcement.

5. The increase in composite fracture strain may reflect a deterioration of the fiber/matrix interface.

6. The glass-reinforced plastics are the only composites that exhibit

increased elongation at 77 K.

7. Tear-type fractures seem to be more evident at low temperature for

epoxy matrix materials.

For glass-fabric-reinforced epoxy composites stressed parallel to the

principal fiber direction, there is a transition in the stress-strain (20, 22,

36, 37) behavior below room temperature that results in both an initial and

secondary modulus characteristic, termed a "knee" of the stress-strain curve.

Figure 1 is a typical illustration of this effect. Initially the knee seems

to be due to interface failure and matrix cracking. There is some evidence to

indicate that the decrease in secondary modulus with decreasing temperature,

as illustrated in Figure 1, reflects a decrease in the shear strength of the

resin matrix or fiber/matrix interface.

530
5. In Summary and Concluding Remarks

For gl ass-fabri c-reinforced nonmetallics this review has demonstrated,

not only a dearth of adequate test and analysis techniques for predicting

strengths and determining mechanisms of deformation and failure, but also, a

need for such analyses and tests.

The "rule of mixtures" was an early and relatively simple model for

relating fiber and matrix strength and stiffness to estimates of composite

strength. Implicit in the model was the assumption that, under an applied

load, there was a constant strain throughout the composite. The interpretation based on that assumption indicated that the first fiber to fracture led

to crack propagation to failure for the composite with fracture occurring at

the failure strain of the fibers.

More recently some investigators (1, 5, 26) have pointed out that the

"rule of mixtures" analysis is inadequate for composites reinforced with very

brittle materials such as glass fibers. Failure for these materials has

statistical features that reflect the nonuniform strength distribution of the

fibers' inherent flaws and imperfections.

Cryogenic testing of these composite materials has been limited to characterization studies which show strength and stiffness increases with decreasing

temperatures. These effects seem to be more a reflection of changes in the

reinforcement than the matrix. There are, however, significant structural

changes going on in the resin at cryogenic temperatures that may or may not be

important in trying to deduce how a glass fabric reinforced epoxy will fail.

Since the fracture strengths of the individual fibers or fiber bundles -

covers a range of values, determination of an appropriate stress for describing

the failure stress of the reinforcement - and ultimately the composite - is a

531
complex problem still to be addressed. Adding to this complexity are important

considerations relative to the sequence of constituent failure events, their

interaction, and contributions to total composite failure.

The initial part of this report has been directed toward a comprehensive

synthesis of behavior and properties information as well as failure models

derived from published investigations of glass-reinforced plastic composites.

The analytical and descriptive material and models included herein have

served as an excellent basis for the low-temperature fractographic analysis of

G-10CR that follows in this report.

C. Low-Temperature Effects on Fracture Modes of G-l OCR Tested

in Compression at 295 K, 76 K, and 4 K

1 . Introduction

To provide an initial data base for engineering applications a cryogenictemperature-characterization program was conducted at the National Bureau of

Standards (1978-79) for the cryogenic-grade gl ass-fabric/epoxy composite,

G-l OCR (37). These experimental studies included electrical, thermal, and

mechanical property characterizations of CR-grade materials over the 295 K to

4 K temperature range.

The objective of the present study, described in this paper, is to assess

the effect of cryogenic temperatures on the failure behavior of the G-l OCR

specimens tested in compression under the program described above. The following specific activities were directed toward achieving this objective:

1. Examination and analysis of fracture surfaces of G-10CR specimens

previously tested in compression at 295 K, 76 K, and 4 K with fractographic methods.

532
2 . Identification and characterization of primary composite failure and

secondary constituent failure as well as the interactive effects of

the constituents that led to gross composite failure.

3. Correlation of failure morphology with room and cryogenic compressive strength characteristics. As a result, the development of an

interpretation of the experimental results, which shows compressive

strength and data variability increasing with decreasing temperature.

4. Comprehensive review of the literature of existing failure analysis

information and models that can be used as a frame of reference for

identifying and elucidating G- 1 OCR compression failure behavior at

room and cryogenic temperatures.

Generically failure was considered to be the occurrence of any discernible

discontinuity in the response of the composite and constituent material to a

mechanically applied external compressive load. Gross composite failure was

determined to be primary, constituent failure to be secondary for comparative

and descriptive purposes rather than to designate order of importance.

Several kinds of failure mechanisms within this classification were used

to identify failure in the G- 1 OCR specimens. Breakdown of the unit structure

leading to failure of the constituents or the bond between them; reduction

in the load-carrying capacity due to localized damage or separation;

significant loss of strength due to crack growth; or total composite

fracture were considered to be identifiable failure processes in this

analysis

.

This study involved an after-the-fact critical examination and analysis

of the fractured specimens. In the analysis it was important to determine:

1. What constituted failure.

2. What was the predominant composite failure mode.

533
3. What were the modes of constituent failure and the probable occurrence of each.

4. What was the probable sequence of failure events in total composite

fail ure.

5. How do the above (i.e., 1-4) correlate with decreasing temperature,

295 K to 4 K.

2 . Material Specifications, G-l OCR

Specimen configuration with adhesively bonded end caps is shown in Figure

2 (37). The end caps were designed for and used to eliminate premature specimen

failure and to assure uniform load transfer as well as good vertical alignment

of specimen ends to avoid column buckling. Figure 2 also shows diagrams of

laminate configurations.

The reinforcement is a silane- finish. E-glass, 7628 woven glass fabric,

plain weave with interlaced warp (length) threads (43/in) and fill (width)

threads (32/in). This type of reinforcement results in a high-volume fraction

of fibers in the resin with a correspondingly high strength and modulus and

different strength properties in directions parallel and perpendicul ar to the

warp.

The matrix is a heat-activated, amine-catalyzed bisphenol A solid-type

epoxy resin, a cross-linked molecular system. Resin content ranges from 32 to

35%.

The composite specimen is made up of layers of similar parallel lamina,

as illustrated in Figure 2. Material for the specimens was obtained from two

producers. Tables 1 and 2 contain typical property values for the E-glass

reinforcement and an epoxy resin (9, 22, 23).

534
3. Test Procedures

Static compression testing was done by applying a compressive load in the

plane of the cloth. The compression test fixture is shown in Figure 3. The

testing environment included air at 295 K, liquid nitrogen at 76 K, and

helium at 4 K. Tests were performed parallel to both warp and fill, and

compressive strength values were determined for both directions at the above

temperatures

.

4 . Fracture Analysis

Fractured specimens of G-10CR tested in compression at 295 K, 76 K, and

4 K were subjected to a critical examination using optical microscope and

scanning electron microscope (SEM) fractography to assess damage and determine

failure modes, both primary and secondary, for each specimen. Failure was

defined in the preceding part of this report. Both macrolevel and microlevel

observations and analysis were considered to be important to the results of

the study. The former assumed the composite to be a homogeneous solid and an

averaging process was utilized to assess damage and failure. The latter was

based on the assumption that the material is a heterogeneous solid and that

constituent (fiber, matrix, and interface) characteristics are integral considerations in composite fracture.

For both primary and secondary failure mode analyses, location and direction of failure and fail ure-surface morphology were stages considered important

to fracture characterization. Determining individual constitutive failure

events as well as their relation to total composite fracture were subsets in

the major focus of the study. Although there seemed to be a significant

degree of consistency in the G-10CR fracture morphology, it still became an

535
extremely difficult task to establish the path of fracture and trace it to

i nitiation

.

In general, damage modes of importance to this analysis were considered

to be cohesive failure of the matrix and reinforcement and adhesive failure of

the fiber/matrix interface.

Constituent damage and failure types, which combine in a variety of ways

as they affect the response of a material to the applied compression load, are

consistent with the numerous published studies summarized in an earlier part

of this report. Those failure modes identified as important to this study

were:

1. Matrix (epoxy) cracking and fracture

2. Reinforcement (glass fabric) microcracking and fracture

3. Fiber/matrix interface, debonding, delamination, and fiber pullout.

Primary compressive failure modes were either (1) an approximate 55°

slant type; (2) a V-notch or chisel type, or (3) a mixed mode or complex type.

Types (1) and (2) have been diagrammed and are illustrated in Figure 4.

Figure 5 correlates each of these three modes [i.e., (a), (b) or (c)] with

temperature for warp and fill directions.

Tables 3 and 4 are temperature-related summaries of significant factors

in the characterization and fracture analysis of these materials. These

tables represent the same information, except that one is a summary for the

warp direction and the other, for the fill direction.

In the first part of the table, failure modes at each testing temperature

have been correlated with the lowest and highest values of compressive strength.

The second part is a summary of the percentages of each primary failure mode

536
for each testing temperature. Several factors became apparent in organizing

the summaries. There is more variability in the primary failure mode in the

warp direction than in the fill direction. In the fill direction, the singleplane ^55° slant-type mode was the predominant mode. This mode represented

100% of the lowest compression strength-vs .-temperature values in the fill

direction, but only 33% of those values in the warp direction. Highest compressive strength-vs. -temperature values were identified as either a singleor a double-slant (shear-type) failure for the fill tests. The warp-test

fracture modes, in this case, were not representative of any highest value

mode except a complex type indicating considerable variability.

The third part of the table identifies the fracture edge for the primary

mode when that mode is either the ^55° slant or V-notch. This edge is a

transverse crack across the specimen with some evidence of tearing at the

lower temperatures (Figs. 4, 6-20).

Part four of the table is a summary of constituent damage or failure-vs.-

temperature. The next part of this report emphasizes in more detail some of

the fractographic analysis that integrates with and documents certain aspects

of the fracture interpretations.

Property measurements determined during the initial characterization

tests are graphically presented for inclusion in this part of the report in

Figures 21 through 23 (35). These diagrams serve as a basis for analysis in

the development of this failure-analysis study. Also, results determined from

these graphs have directed construction of additional graphical comparisons

considered to be significant and illustrated in Figures 24 through 26.

Figure 1, as previously indicated, illustrates a typical tensile stressstrain diagram for some of the G-10CR specimens tested in the warp direction

537
at room temperature and 76 K (37). The diagram, even though it represents a

typical tensile property, has been included with the report because it is

representative of the stress-strain behavior for this material. The work

required to fracture a specimen can be related to the area under the stressstrain curve and is an indication of the toughness of a material. Using this

area as a measure, the work at fracture shows an increase when the temperature

is lowered to 76 K. At 76 K there is a discontinuity in the configuration of

the curve, considered to be a "knee" at about 300 MPa with an accompanying

change in the modulus. Initial and secondary moduli reflect this transition at low temperatures but not at room temperature. This type of behavior

can be correlated with specific failure occurrences i.e., beginning matrix

cracking at the knee and a degradation in interface shear strength reflected

by the lower secondary modulus.

Figure 21 has been included to provide a graphic comparison of the range

of measured test values for compressive strength. Young's modulus, and tensile

strain to failure for room temperature, 76 K, and 4 K in both warp and fill

directions. The number of specimens for each test condition are shown in

Tables 3 and 4. The bars at each temperature in Figure 21 are representative

of the highest and lowest test value and thus illustrate the variability in

the data for each of the properties shown. Figure 22 diagrams all measured

values for Young's modulus vs. temperature and Figure 23 is a graph of all the

measured data for compressive strength vs. temperature. The outer straight

lines have no physical meaning other than to indicate upper and lower bounds

for the data and to present a graphical comparison of the scatter in the data

at each temperature; the center straight line connects the average value at

538
each temperature. The temperature-related vertical line delineates the

calculated standard deviation for the range of values.

It's interesting to note the significant increase in strengths for warp

and fill directions compared with those at room temperature. Figure 24 illustrates a graphical comparison for warp and fill tests of the low temperatureroom temperature percentage increase in compressive strength. Young's modulus,

and tensile strain to failure.

Another interesting factor to be noted in Figure 23 is that the low

temperature variability (i.e., the scatter in the data) is reversed for

measured compressive strength values in the warp and fill directions. At 76 K

the scatter is very large for the warp direction and is significantly less in

the fill direction. At 4 K the reverse is true. Variability (as expressed in

the standard deviation and coefficient of variation for compressive strength,

Young's modulus, and tensile strain to failure) has been addressed in Figures

25 and 26. These graphs summarize and compare variability as a function of

temperature

.

The graph illustrating Young's modulus vs. temperature (Figure 22) was

determined and presented in a manner similar to the one for compressive strength.

Earlier studies have indicated that the low-temperature related increase in

composite strength and modulus primarily reflect similar changes in the glass

reinforcement. Plotted values for Young's modulus in warp and fill directions

indicate, as was also illustrated for compressive strength, that there is

significantly greater scatter for the data in the warp direction. Variability

at 4 K does not appear to be a significant factor for either warp or fill

directions

.

539
Optical and scanning electron microscopes were used to examine details of

the fracture surfaces to identify and evaluate constituent failure. Figures

6 through 15 are typical photographs obtained with the optical microscope.

Figures 16 through 20 are typical SEM photographs of fracture details for

specimens tested in the warp direction at various temperatures . Significant

details have been identified and delineated on each of the photographs. These

illustrations can, in turn, be correlated with the summaries in Tables 3 and 4.

5 . Results and Discussion

(a) General

As indicated in an earlier part of this report, previous studies have

shown that the response of a glass-reinforced laminated nonmetallic tested in

static compression is a function of the type, orientation, and volume fraction

of the reinforcing material; matrix characteristics; precision of alignment of

the test specimen; and direction of loading (e.g., load-path distribution as a

function of the anisotropy of the material). A woven glass-reinforced epoxymatrix composite subjected to an externally applied compressive load parallel

to the fibers (either warp or fill) can undergo in-plane shear deformation

(affecting both the lamina and the interlaminer areas), longitudinal and

transverse bending of the reinforcement, and possible bending of the specimen.

The matrix undergoes shear deformation. Bending of the fibers results in

secondary tensile and shear stresses that may be highly significant to the

failure mechanism.

Published studies on the compressive strength of laminates have identified shear and interface (debonding, delamination, and pullout) failure and

fiber buckling as primary failure modes. The shear mode, however, has seemed

540
to predominate. A propagating shear mode of failure across a specimen can be

the result of compression-induced bending or shear stresses (4, 5, 28, 31,

33).

There are several constituent structurally related responses to the lowtemperature environment that are significant to the behavior of glass-rei nforced

nonmetallics. Failure modes of the composite are in some way related to

interacting modecular and microstructural (constituent) events. An increased

rigidity in the polymer chain degrades its ability to dissipate energy. Chemical

structure, in terms of cross-linking and directional characteristics inherent

in primary covalent bonds and secondary Van der Waals forces, characterizes a

molecular anisotropy in the matrix resin, which can be significant to composite

behavior. Hartwig has shown that for an epoxy resin: (1) Young's modulus and

Poisson's ratio show little or no dependence on chemical structure at low

temperature; (2) tensile strength shows some dependence on cross-link distance

at temperatures below 76 K; and (3) the energy of fracture shows significant

structure sensitivity at both normal and cryogenic temperatures. The energy

of fracture is released locally and can lead to a significant temperature rise

in the specimen. Hartwig points out that epoxy resins behave viscoelastically

at or near the transition temperature and linearly elastic to fracture with an

accompanying low fracture strain (approximately 2%) at cryogenic temperatures

.

Increased strength, brittleness, and resin rigidity characterize the matrix at

the low temperatures (15, 35). Increased strength and fracture strain are the

cryogenic effects on the glass reinforcement.

The topography of a fracture surface is largely a function of the stress

conditions at fracture. An analogy with metals would indicate that lower

stress levels are generally identified with a lower level of energy being

dissipated at fracture and a smooth, regular fracture surface.

541
(b) Fractographic Observations and Analyses

With the above as an introduction, the following relates to the specific

compression tests of G-10CR specimens at 295 K, 76 K, and 4 K and the fractographic examination and analysis of those specimens that comprised this study.

There were two significantly different testing conditions for the G-10CR

specimens compression tested at each of the indicated temperatures. Specimens

were oriented with the compression load parallel to warp threads for one

series of tests and to the fill threads for the other group. (There was a

third test orientation normal to the fabric reinforcement. These specimens

however, were not included in this fractographic assessment.) The ratio of

fill-to-warp threads is approximately 0.75. Measured values of Young's modulus,

c u E, and compressive strength, a , for the various testing temperatures correlated well with this ratio.

As the summary in Tables 3 and 4 indicates a single-plane, slant-type

primary fracture configuration occurred most frequently for both warp and fill

tests at all temperatures but more frequently for the fill tests. This investigation has identified that mode as an approximate 55° slant-type mode (^55°),

see Figure 4. A second slant-type configuration occurred and was identified

as two intersecting planes showing a "V" or chisel appearance. This configuration was designated as either a V-notch or chisel type (Figure 4). There

were more V-notch failures for the warp tests than the fill tests. For the

former, all occurred at 295 K and 76 K. At 295 K these failures generally

occurred (1) within the gage section but near the face of one of the grips and

(2) with considerable longitudinal delamination. At 76 K the fractures occurred

(1 ) at the face of the grip, (2) with delamination, and (3) with considerable

general damage and surface roughness.

542
There were more complex or mixed-mode fractures in the warp direction

than in the fill direction and more evidence for this at 4 K. There was

significantly more variability in the primary fracture mode for tests performed in the warp direction than in the fill direction, and this variability

increased at the lower temperatures. High values of compressive strength for

warp tests generally correlated with complex fracture modes at cryogenic

temperatures; low values of compressive strength for these tests correlated

with either ^55° slant type or V-notch at 295 K and 76 K and was mixed at 4 K.

For the fill tests at all temperatures , the lowest values of compressive

strength correlated 100% with the ^55° slant mode. Whereas high values of

compressive strength at all temperatures correlated with either the %55° slant

or V-notch mode.

For all tests, both warp and fill, damage and surface roughness generally

increased at the lower temperatures. This damage was manifest in matrix

cracking and greater transverse fiber damage. Individual fiber and fiberbundle fracture, debonding, and pullout were character!' stic of the latter.

Increased damage and surface roughness at low temperatures has been correlated

with energy or work to fracture at 76 K, a measure of which is the area under

the stress-strain curve in Figure 1. Published studies have indicated increased work to fracture is a function of the interface and its ability to

deflect or blunt a crack propagating normal to the reinforcement (4, 5, 12,

17, 22, 27). For the present study, the low-temperature significance of this

effect, in part, can be associated with changes taking place at the interface,

which probably intensify the mechanism for occurrence.

Another characteristic that has been identified with primary fracture in

both warp and fill directions for slant-type failures at all temperatures is a

543
major transverse crack across the specimen, which has been determined to be

the primary fracture edge. This crack, illustrated in Figure 4, shows evidence of being initially a transverse crack in the matrix of the outer lamina

(Figures 6-10 and 14). Fractography shows this edge is deflected very slightly

at fiber interfaces (for those fibers parallel to the applied load) but reinitiates in a transverse direction, propagating and deflecting intermittently

across the specimen face. Reinitiation and propagation are such that the

pattern of the fracture edge is primarily at right angles to the applied load.

At cryogenic temperatures the fracture edge is similarly transverse but there

is evidence of tearing (Figures 8 and 13).

Several factors complicate fractographic analysis in general and the

determination of an initiating failure mode in particular. Composite anisotropy contributes significantly to variability at the interfaces. As this

study shows, there has been more interface damage variability transverse to

the applied load than there has been parallel to it. In conjunction with this

effect is the fact that load-path and stress distributions are a function of

the inherent anisotropy of the composite material. Moreover, the determination of an initiating failure mode is complicated by the differential stiffness,

expansion and contraction characteristics, and subsequent residual curing

stresses developed in fabrication. And, the load-bearing capabilities of

matrix and reinforcement, although integral in forming the unit structure of

the composite, represent a significantly important differential in the loadsupport effect. Thus, the reinforcement carries more of the load than the

matrix because of its higher strength and modulus. The high volume fraction

of the glass reinforcement represents another parameter that makes a significant contribution to the failure mode in compression.

544
Constituent behavior, properties, and interactions have received equal

emphasis in this fractographic study to describe G-10CR fracture behavior. In

examining the fracture surfaces, the analysis was based first on discrete

damage events (i.e., matrix, fiber, and interface failure) and second on the

constituent interaction, which determined a most probable sequence of failure

events. The former was developed in considerable detail and has been typically

illustrated in Figures 12 through 20. The latter is more complex and requires

detailed techniques of analysis, which will direct efforts in future studies.

It has received initial coverage here.

(c) Conclusions

When glass-fabric-reinforced systems such as the G-10CR materials are

subjected to an externally applied compressive load parallel to the fibers at

either room or cryogenic temperatures, the stiff epoxy resin holds the fibers

firmly fixed, constraining deformation which leads to strong stress and strain

concentrations (5, 31, 34). As in other composites under an applied stress,

different displacements occur in the matrix resin and glass reinforcement

owing to an inequality in the moduli (18). The stiffness differential results

in shear strains and stresses, a primary mechanism for distributing the load

between matrix and reinforcement. Because the matrix will deform more than

the fibers and because of the difference in elastic moduli, the latter support

most of the load at all levels. Previous studies (3-15, 18) have shown that

an axial load, parallel to the fibers, is transmitted directly to the reinforcement for continuous-filament composites and via shear stresses at the

interface for discontinuous-fiber laminates. The epoxy matrix acts primarily

as a mechanism for load transfer and stress redistri bution . Although the

545
G-10CR materials are considered to be continuous, in actuality, as with other

continuous gl ass-f i ber-rei nforced composites, continuity characteristics can

be disrupted by fabrication flaws or early fiber fracture at low levels of

applied stress. The latter is the result of the lack of uniformity in glass

strengths owing to defects inherent in all very brittle materials. The shortened fibers, the weaker ones that fracture at low applied-stress levels,

characterize a discontinuous effect in what is normally considered to be a

continuous-filament-reinforced material. Stress perturbations and stress

concentrations ultimately result in weakened zones in the composite that lead

to total composite fracture.

An increase in strength and strain properties of the G-10CR specimens

with decreasing temperatures seems to be more a reflection of changes taking

place in the reinforcement (i.e., an increase in strength and modulus) than in

the matrix.

A fractographic assessment of the specimens has shown some rather specific

consistencies as well as differences in the appearance of the glass fabric

with decreasing temperature. First and most general, the evidence indicates

that the reinforcment fractures by tension or bending. Second and more specifically, there is a brush-type fracture of fiber bundles at the crossovers and

parallel to the applied compressive load for both warp and fill tests (with

little or no debonding of individual fibers); see Figures 14 through 20.

There are more of these fractures visible on the fracture surface at the

higher temperatures than at 4 K. Alternatively, there are less visible brush

fractures at low temperatures for tests performed in the fill direction. A

significant number of the brush fractures are accompanied with delamination

outermost of the fiber crossover area (Figure 24-26). To a lesser extent,

some of the bundle fractures remain embedded (little or no delamination) with

546
only the brush ends showing (Figures 22 and 24). This effect is apparent at

low temperatures when there is a transverse matrix crack contiguous to a

transverse fiber bundle. For both warp and fill tests, fractographic evidence

points to an elastic bending of the fibers with resulting fracture at the

crossovers for those fibers parallel to the applied load. There is a characteristic consistency to this type of fracture mode at all temperatures. The

only temperature-related difference seems to be that fewer of these are visible

on the fracture surface at the lower temperatures, particularly for the fill

tests. A low-temperature-related effect is that more matrix shows on the

fracture surface at 4 K. The fractography has provided strong evidence that

longitudinal bending fracture (i.e., for fibers parallel to the applied load)

occurs as a result of stress concentrations at the fiber crossovers, where the

dissipative mechanism with the highest probability for relieving the stress

buildup is the bending fracture of the bundles. Published fracture studies

reviewed for this report have shown no evidence of the brush fracture mechanism observed in the G-10CR specimens.

For both warp and fill tests, examination of the fracture surfaces of the

specimens showed a significantly greater degree of randomness for damage and

fracture (individual fiber and fiber bundle, debond, and pullout) transverse

to the applied load, at cryogenic temperatures in particular. Room-temperature random effects were significantly less. At room temperature, evidence of

shear failure was apparent for transverse fibers in that damage was primarily

limited to delamination with little or no damage to the fibers in the shear

direction. The undulating nature of the transverse fibers was not destroyed

at the fracture edge (Figures 16 and 17).

547
Finally, as in other published studies, fiber cracking and fracture

appears to be due, in part, to cracks at the surface as well as to intrinsic

defects. As the fibers fractured in the composite, the increasing load was

redistributed by the matrix via the interface to the shortened, weaker as well

as the remaining, stronger, long fibers. Initial fractographic analysis

indicates that this process continued until a significant number of fibers

were reduced to ineffective lengths, causing a weakened region in the composite which, in part, led to total composite fracture.

Matrix fracture was characteristically either tensile or shear, with the

latter predominating. Since the matrix resins serve primarily as a mechanism

for load transfer and stress redistribution, there are certain temperature-related

factors that directly or indirectly determine the effectiveness of those

functions. On the molecular level, for example, the energy dissipative mechanisms

inherent in the mobility of the polymer chain were more constrained with

decreasing temperatures. Changes that occurred in the resin at low temperatures

affected the load transfer function, a parameter that is sensitive to residual

stresses developed during fabrication, to resin brittleness, and to a change

in resin modulus. The ductility of the resin matrix decreased rapidly with

cryogenic cooling. Resin strength and modulus increased; strain capability

was reduced; toughness decreased; and, stress-strain behavior changed from

viscoelastic to linearly elastic at 4 K. At cryogenic temperatures there was

evidence that the epoxy resin matrix of the G-10CR materials cracked before

the glass reinforcement (Figures 1, 9b, 10b).

Although there has been extensive consideration in the published literature about which constituent (the matrix or the reinforcement) , fails first,

this study indicates that the most probable first failure event for these

548
testing conditions was the fracture of the outer layer of the epoxy matrix.

As the matrix crack in the outer lamina developed and propagated transverse to

the applied load and across the specimen, a less effective distribution of the

load resulted. Fibers contiguous to the matrix crack became more highly

stressed than those firmly embedded, and some bending of the fibers occurred

in the vicinity of the cracked matrix. This process continued as the load increased and contiguous regions were weakened. Several factors lend support

to matrix cracking preceding significant fiber fracture: First, epoxy strength

is lower than that of the glass reinforcement. Second, the nonuniform strengths

characteristic of glass fibers signify a randomness in their total fracture

pattern for which there is little or no evidence in the constituent failure

modes parallel to the applied load at any temperature. There was, however, a

distinct randomness in failures transverse to the applied load with a more

irregular effect becoming apparent at the lower temperatures. Finally, typical

tensile stress-strain curves at 295 K and 76 K for these specimens show a

transition in the configuration occurring at 76 K, a "knee" with an accompanying

secondary modulus (Figure 1).' The appearance of a "knee" has been correlated

with initial matrix cracking and the secondary modulus, with interfacial failure.

There were several fractographic features that provided evidence of

changing matrix damage at the lower temperatures. The increased brittleness

of the matrix resulted in opaqueness. Tear features were apparent at the

lamina interface and tear ridges showed in the matrix (Figure 19). A number

of matrix cracks, both external and internal were visible. There was more

matrix on the fracture surface for the low temperature specimens and there was

some evidence of hacklelike marks of the matrix at the fracture surface

( 21 ).

549
Finally, interface failures at all temperatures for these specimens were

identified as delamination (of the lamina) and debonding and pullout of the

fibers and fiber bundles. Debonding and pullout were most characteristic of

failures transverse to the applied load. There was significantly more of

this type of damage at the lower temperatures. Delamination was identified in

several ways. First, there were the short delamination steps associated with

brush fracture of the longitudinal bundles (Figure 17). Second, there was

almost total delamination of the lamina in a direction parallel to the applied

load for some of the mixed mode failures. Third, there was the transverse

delamination at the primary fracture edge for the warp tests at 295 K, which

left the undulating surface of the fill bundles intact (Figure 17).

There is some evidence in the published literature to indicate that an

increased fiber fracture strength or a decreased matrix or interface shear

strength lead to debonding rather than fracture in the debond/fracture subsequence of events. There seems to be some support for this sequence at the

lower temperatures of the present study.

In conclusion, examination and analysis of the fracture surfaces of the

G- 1 OCR specimens fractured in compression at 295 K, 76 K and 4 K have provided

a basis for the identification and characterization of primary composite

failure and secondary constituent failure. As the summaries in Tables 3 and 4

indicate there are identifiable tempera ture-rel ated failure effects. Figures

6-20 correlated with Tables 3 and 4 emphasize the major effects. Figures 21-

26 illustrate temperature related increases in strength stiffness, and variability or scatter in the data (in terms of standard deviation and coefficient

of variation). This information is also reflected in Tables 3 and 4 and can

be correlated with Figures 6-20.

550
6. Recommendations for Future Studies

There has been little or no fractographic or modeling studies of the lowtemperature compression failure behavior for gl ass-fabric-rei nforced epoxies.

The in-depth review of published studies that comprised the initial part of

this report has been synthesized to provide a comprehensive survey of relevant

information that can be used to support and direct the investigative efforts

of this study. The literature review has been immeasurably helpful in the

development and organization of this report and the failure studies that

preceded it. The explanations, interpretations, and extrapolations have

provided an invaluable documentary information base.

A logical follow-up of this temperature fracture-mode study should

include reconstruction of the fracture process from initiation, through propagation, and on to final fracture developed in terms of a most probable

sequence of failure events by using a Fault-free Analysis, which has been

described in other publications (10).

Concurrent with a sequence-of-events analysis should be a stress analysis

describing material behavior just prior to fracture. The fractography and

modeling developed in the present report will provide support information for

the analyses suggested above.

These stages, in turn, can serve as a logical supplement to the systematic development of a theory of failure for analyzing the low-temperatuare

behavior of glass-fabric-reinforced nonmetallics under a static compression

load.

551
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556
List of Tables

Table 1. Typical Values,

Table 2. Typ

Tabl e 3. G-l

Table 4. G-l

ical Values,

OCR Fracture

OCR Fracture

E-glass

General Purpose Epoxy Resin

Analysis Summary, Warp Tests (Compression)

Analysis Summary, Fill Tests (Compression)

557
Table 1. Typical Values, E-Glass (9, 22, 33)

E-glass

Properti es Temperature

295 K

~ 230 K~ 205 K ~83 K

Tens . St. , MPa 3447 5171 5343 5654

Young's Modulus,

7? 4 Modulus increases approximately 3%

E, GPa between 295 K and 76 K.

% Fracture Elongation ~ 2 . 5%

Table 2. Typical Values, General Purpose Epoxy Resin

Epoxy

Properti es Temperature

295 K 77 K 20 K

Tens . St. , MPa 41 103 100

Compress . St. , MPa 90-138

Young's Modulus, E, GPa 21 47 66

% Fracture Elongation 4.5% n 1 . 8%

558
Table 3 G- 1 OCR Fracture Analysis Summary, Warp Tests (Compression)

295 K (8 specimens) 76 K (7 specimens) 4 K (3 specimens)

cu

0 vs. Failure Mode:*

c u o vs. Fail ure Mode:*

c u a vs. Failure Mode:*

(a) low aCU: '55° slant (a) low o CU: V-notch (chisel) (a) low a CU: mixed

(Failed at grip--no gage section

separation)

(b) high o CU

: V-notch (chisel) (b) high a CU

: mixed mode (b) high a CU

: mixed (55° slant,

delamination, and bending.)

Primary Mode: Primary Mode: Primary Mode:

37%

25%

38%

'55° slant V-notch

Miscellaneous (broke in gage section, mixed or missi ng)

14% '55° slant 43% V-notch at face of grip

43% other (mixed and saw-tooth)

67% mixed

33% failed at grip--no gage section separation

Primary Fracture Edqe Primary Fracture Edqe Primary Fracture Edqe

Major transverse crack

across specimen.

Initially cracked through

matrix of outer lamina.

Major transverse crack across specimen; same as 295 K. Some tearing-type damage at edge.

2 specimens show major transverse

crack similar to 295 K and 76 K. No gage-section separation of third specimen.

Constituent Failure Constituent Failure Constituent Failure

(a) Longitudinal Direction (warp) Brush fracture of fiber bundles with some slight

delamination at fabric crossovers

More general damage and more random

in nature than 295 K. Rough fracture surface

(a) Longitudinal Direction (warp) some brush fracture and delamination of fiber bundles

Considerable damage at primary transverse crack.

(a) Longitudinal Direction (warp) some brush fracture longitudinal fiber bundles with some microcracks in fibers and delamination.

(b) Transverse Direction (fill) some fiber microcracks, breaks

and pullout.

Many fiber bundles show delaminated

outer surface but imbedded in inner matrix resulting in an undulating

surface.

(b) Transverse Direction (fill)

Considerable damage to transverse

bundles and individual fibers fracture, debond, and pullout. Bundles pulled out--not imbedded

in inner matrix as they were at 295 K.

(b). Transverse Direction (fill)

microcracks in fibers. Some debonding, fiber fracture

and pul lout.

(c) Interface Damage

Primarily delamination

(c) Interface Damage

Debonding, delamination

and pullout

(c) Interface Damage

Considerable delamination

some fiber debonding

(d) Matrix (d) Matrix (d) Matrix

Glassy and transparent.

Smooth at fracture edge.

No evidence of matrix cracks outside the fractured area

Opaque. Microcracks in gage section on reverse side of fracture surface.Opaque. Si gni ficant matrix

damage--cracking and roughness. More matrix showing on fracture surface than at higher

temperature. Cracks showing

on external specimen surface

at warp and fill crossovers.

*Note: c CU refers to ultimate compressive strength.

559
Table 4. G-10CR Fracture Analysis Summary, Fill Tests (Compression

295 K (3 specimens) 76 K (3 specimens) 4 K (3 specimens)

o CUvs. Failure Mode:

c u o vs. Failure Mode: a CU vs. Failure Mode:

(a) low a CU

: '55° slant (a) low a cu

: '55° slant (a) low oCU

: '55° slant

(b) high o CU

: '55° slant (b) high a CU

: V-notch (b) high o CU

: '55° slant

Primary Mode: Primary Mode: Primary Mode:

67% '55° slant

33% Missing Spec

.

67% '55° slant

33% V-notch at grip 100% '55° slant

Primary Fracture Edge Primary Fracture Edqe Primary Fracture Edqe

Major transverse crack across Major transverse crack across Major transverse across specimen

Initially cracked through

matrix of outer lamina

specimen; see 295 K. More tearingtype damage

specimen; see 295 K.

Constituent Failure Constituent Failure

(more surface damage than at 295 K) Constituent Failure:

(Not as much visible fiber damaqe

on fracture face; more matrix J showing compared to that at 295 K

(a) Longitudinal Direction (fill) (a) Longitudinal Direction (fill) (a) Longitudinal Direction (fill Brush fracture with some Brush fracture of fiber bundles Few bundle fractures delamination at crossovers with more delamination than room with some delaminationof fiber bundles. temperature. Microcracks in fibers. mostly at fracture edge.

(b) Transverse Direction (warp) (b) Transverse Direction (Warp) (b) Transverse Direction (warp) More random fiber More transverse fiber bundle Considerable fiber damage;

damage than warp tests. Con- debonding, fracture and pullout. irregular and random siderable damage of fibers Less individual fiber debonding damage of individual debond, fracture and pullout. fracture and pullout. Considerable transverse damage at fracture edge

.

fibers and bundles;debonding

fracture and pullout.

(c) Interface Damage (c) Interface Damage (c) Interface Damage

Delamination. Debonding and pullout of bundles and individual

fi bers.

Delamination, debond, and pullout Some longitudinal delamination. Transverse fiber

and bundle debond and pullout.

(d) Matrix (d) Matrix (d) Matrix

Glassy and translucent Opaque. More broken up and cracked

than at 295 K. Significant matrix cracks undercutting transverse fibers. Trans- verse crack of matrix on reverse side of fracture surface.

Opaque. More matrix

showing on fracture surface

Small number of glassy

regions. Irregular surface

Some internal cracks

showing on reverse side.

560
LIST OF FIGURES

Figure 1 .

Figure 2.

Figure 3.

Figure 4.

Figure 5.

Figure 6.

Figure 7.

Figure 8.

Figure 9.

Figure 10.

Figure 11 .

Figure 12.

Figure 13.

Figure 14.

Typical tensile stress-strain diagram G-10CR, Warp (37).

Laminate and specimen configurations used in determining

static compressive properties of G-10CR at 295 K, 76 K,

4 K.

Fixture for static compression testing at cryogenic temperatures (13).

Typical primary fracture modes in compression. G-10CR at

295 K, 76 K, 4 K.

G-10CR. Primary compression fracture mode vs. temperature.

See Tables 3 and 4 for information on the number of specimens

in each machi ne mode.

G-10CR, warp test, 295 K, ^55° slant mode fracture (37).

G-10CR, fill test, 295 K. ^55° slant mode fracture (37).

G-10CR, warp test, 76 K. (a) ^55° slant mode (b) and

(c) mixed-mode fracture, (a), (b), and (c) are different

specimens (37):

G-10CR, warp test, 76 K. (a), (b), and (c) are different

specimens (37).

G-10CR, fill test, 76 K. (a) and (b) are different specimens (37).

G-10CR, warp test, 4 K mixed mode failure (37).

G-10CR, warp test, 4 K. (a) and (b) are different examples

of mixed mode failure (37).

G-10CR, fill test, 4 K. ^55° slant mode failure (37).

G-10CR, warp test, 295 K. (a) and (b) are different specimens .

561
Figure 15.

Figure 16.

Figure 17.

Figure 18.

Figure 19.

Figure 20.

Figure 21 .

Figure 22.

Figure 23.

Figure 24.

Figure 25.

Figure 26.

6-10CR, warp test, 295 K. Fracture surface (35).

G-10CR, warp test, 295 K. SEM of fracture surface (37).

G-10CR, warp test, 295 K. SEM of fracture surface (37).

G-10CR, warp test, 76 K. SEM of fracture surface (37).

G-10CR, warp test, 4 K. SEM of fracture surface (37).

G-10CR, warp test, 4 K. SEM of fracture surface (37).

G-l OCR . Compressive strength. Young's modulus, tensile

strain to failure vs. temperature (37).

G-10CR. Young's modulus vs. temperature (37).

G-l OCR. Compressive strength vs. temperature (37).

G-10CR. Percent increase in tensile strain to failure.

Young's modulus, and compressive strength at cryogenic

temperatures relative to room temperature. Calculated

from results in reference 37.

G-10CR. Standard deviation vs. temperature (37).

G-l OCR . Coefficient of variation in percent vs. temperature (37).

562
STRESS,

<r,

MPa

Figure 1. Typical tensile stress-strain diagram, G-10CR, warp (37).

563
Fill 2.75

(7.0)

atlD C

0

I TOP

1 ypicai

Fiber

Crossover

Warp fiber bundles

Fill fiber bundles

(a) FRONT

Diagram showing top and

front views of woven fabric

(with fiil and warp bundles)

(b)

Diagram of warp and

fill plain weave in a G-10 compression specimen.

(c)

G“ 1 0Cr

compression specimen

Specimen

0.15x0.3

(.38x.76)

Sec A-A

Parallel layers

of lamina

4.7

( 1

1

. 9 )

Delamination surface

showing fiber weave

/

0.75

(1.9)

r

1.0

(2.5)

i

/6061-T6

Al grip

Al spacer

2.75

A^(7.0)

(d)

G-IOCr compression specimen

with aluminum grips.

Units: in. (cm) (37)

Figure 2. Laminate and speciman configurations used in determining static

compressive properties of G-IOCR at 295K, 76K and 4K.

564
Figure 3 Fixture for static compression testing

at cryogenic temperatures (13).

565
Typical fiber

crossover

Primary fracture edge

Matrix missing.

Exposed fabric

weave showing

(a) Primary compositefracture ^ 55® slant

Matrix missing.

Exposed fabric

weave showing

Primary fracture edge

(b) Primary composite fracture,

V-notch or chisel

Matrix missing.

Exposed fabric

weave showing.

Primary fracture edge

Major transverse crack

across specimen.

Initially cracked through

matrix of outer lamina

Primary fracture edge

(c) Primary ~55® slant fracture surface

Figure 4. Typical primary fracture modes in compression.

G-l OCR at 295K , 76K, 4K.

566
PRIMARY

FRACTURE

MODE,

%

100

80 Warp Test

100

80

60

40 1-

20

0

55° slant

Fill Test

*55° slant

Missing

specimen^

V-notch (chisel)

4 K 76 K 295 K

Figure 5. G-10CR. Primary compression fracture mode vs. temperature. See

Tables 3 and 4 for information on the number of specimens in each fracture

mode.

567
<uO CD

ra on 4- s- 5- <u 3 >

on onc

CD ro

S- S- 3 4->

4-> o Cl) fO '

L_ 4-> 4— 4->

CD r—

SZ

+-> >> 1

4- CD r-v O > CO

•r— '—

x 5 +>

CD ro CDS- > CD 3

S- 4~> XJ O

CD cn rO cn d CD S- S- -r- cn 4—

ra 5 ro r— O E CD C _d ro “O

LxJ cn “O

OE

-O +-> c

rO

on o

LDLO

LO

cn

c\j

O) ~ao raCl c

ra r— r-* ui lo

ro o

LO II LO 3o

<? o

CDC

4-> rO3

T3c3

cn

u c -I-

•«- s- ^ -a O ro _d 4-

00

CD ^ _c

rO -+-3

L|

—O o u S -r- CD 4~> •r- on > -r-S- TD <D

CD +-> CJI O

S- ro

rO i- <— rO C -C

LxJO

CL

irO

ccooo

LO

CD

s~3

CD

LU

b68
569

Figure

7.

G-10CR,

fill

test,

295

K.

^55°

slant-mode

fracture

(37)
>

CD

(D CD £ cr» O -O

GO CD

CD rC • S_ CL

CD 3 21

T3 +-> OO^f- E <T5 S_ O 4-> 4- O'* c

(T3 >*> r~~ S— I! CO 03E 13 O *r- o

LO S- LO CL

O ?

f0

03 o

Csj COCO

II 3U o

CD

i-3

CD “OO

s:“O

CDX

raCL

CO

r^

CD X3O

s:

"O

CDX

570

Figure

8.

G-10

CR,

warp

test,

76K.

(a)

^55°

slant-mode

(b)

and

(c)

mixed-mode

fracture
X

*i

—

4-> 03 CD Z cn

cn c

~o CD *r~

C CD CO j*: C a

•r— CD CD rO

4-> >•> s~ > S- CJ CL S- 3 to O

4~> CD S- rO 4-> c

CO ro E O rO CD •r- ^ •r- 03 S4-> -a '

S^- S- a_ (4- 1—

<uCD

rocnU

a>Cl

IS) 4-> U

CD CL

5*-

*«— 03 O 3:

CO ro 00 CL o 2:

S-u CO

ro

CXJ cn co c coI o

03=3 s-u o

s~

4->

CO

CL

S03

CDU

ro 4-

S~3

oo

CD

S~3

4~> u

rO

SCOo

CD

i-3

CD TOO

~a

CDX

roCl Zocn

CL

roO

"Oc

CD

CL 3

>•»O

CD

S~

roO

•oC

03

•O

LO

CD

4-> CL

ro

SC_> o

CD3

cn

different

specimens.

(37)
test direction — (fill)

primary fracture edge

transverse (warp)

direction

(a) ^55° slant mode. Considerable transverse

(fill) damage. j

CU = 536.6 Mpa

(b) Matrix cracking across specimen gage

section face a = 561 .6 Mpa

Figure 10. G-10CR, Fill Test, 76K (a) and (b) are different specimens. (37).

572
CD

O

~o

U X) .— c

03

O c •*

•r— -M “O 03 3 d 4_> •»

1 E

cn 03 d r— O CD _J “O

03d

“O +-> M- 3 00 C

4-> <D

•r- I— 00 C3 *r— d X O • 03 r CDC O CO * -r- 4-> CL

4-> -XD IEI e c: r—

03 CL) QJ of) r- Dr- • oo «— ^3- X5 CO ID o c: 5- co LO C0 (t II LO QZS * U * d 00 D •• O -r- <D -r- “O 4-> C o «0 O 00 E d -r- CL

•p— 4-> 03 "O E u o

<D 03 CD X r— d -D

•r- CD *rC

21 “O “O O)

03

CL

03O

~od

O)

CD -3O

M— E o +-> CD d O 03

03

Li_ CO —

d

03

O

d t-

•r- +-> T3 03 3 d 4-> -p-E

cn 03 C r— O <D —J “O

573

(c)

Enlarged

view

of

fracture

at

right

end

cap
(a) Mixed mode failure: delamination, bending

of the specimen, and slant mode fracture.

a

cu = 881.9 Mpa

(b) Mixed Mode Failure. No gage section

separation. o CU

- 849.9 Mpa

Figure 12. G-10CR, Warp Test, 4K. (a) and (b) are different

examples of mixed mode failure (37).

574
warp (

transverse)

di

recti

on

S-Q.

O

O *i— LO

4~> 03 I! C

•r-3 EU

03

O

CL) • “O CDO

r— 03 03 M— e S-

•r- Z3 "O oo 3

•r- d) CD E C *r— O o

r— (D

Q_

JZ 00

<D “O

O

O 4-> cc

+-> Cl) C r—

03 •

r— 03 CO S~03 O CL LO

LO

>> S- S- 3

03 4-> E O CD T- 03 cr> S- s- -O

Cl <4- CD

co

CD3

CD

S~3

C3

(c)

Enlarged

view

of

the

fracture

Surface

of

another

specimen, a=

593.3

Mpa
delamination

fracture face

Fill

Warp

undulating

geometry of woven fabric

individual lamina

fill (transverse) fiber bundle

fracture and pullout.

Crush

fracture warp

( 1 ongi tudi nal

)

fiber bundles

fill

(b) primary (transverse) fracture edge.

Figure 14. G-10CR, warp test, 295 K.

(a) and (b) are different specimens.

(37)

576
fill fiber

bundle debond,

fracture, and

pul lout

.

brush fracture

of warp fiber

bundles - with

some delamination

primary

( transverse)

fracture edge

Figure 15. G-10CR, warp test, 295K.

Fracture surface ( 35)

.

Warp

(test) direction

typical brush

fracture of

v^arp fiber bundles.

Fractured fiber

ends are showing.

i i 1 1 ( transverse'

undulating

fill fibers

warp

(test direction)

Figure 16 G-10CR, warp test, 295 K. SEM of

fracture surface (37)

577
primary

( trans verse)

fracture edge

primary

( transverse)

fracture edge

matrix

(a)

fill (transverse)

Matrix steps. Delaminated regions showing missing warp bundles.

End views of fractured warp bundles are visible at rear boundaries of matrix steps.

figure 17 s- 1 OCR , warp test, 395K. SEM of fracture surface (37).

578
fill fiber

debond, fracture,

and pullout

PWrwBfc i

(a)

matri

warp fiber

fracture and pullout

fracture edge (transverse)

Figure 13 G- 1 OCR , warp test, 76K. SEM of fracture surface (37).

579
fill fibers

end view of

fractured warp

fiber bundles brush

fracture

fill (transverse)

Primary fracture edge

matrix cleavage marks

warp fiber

bundle brush

fracture

fill (transverse,

warp (test

di rection

(b)

f iqu re 1 9 G - 1 Of,

1 wa rp te s t , 4K

Sff' of fracture surface (37)

580
matr.i x

C 1 savage

marks

fill fiber debond,

fracture and pullout.

warp fiber bundle

brush fracture - with

delamination

warp (test direction^

fiber crossover

Brush fracture warp fiber

bundles. Some delamination.

Inclined fracture surface

of specimen. Shows undulating

geometry of fill fiber bundles

Some fiber fracture and pullout

Delaminated lamina,

showing woven glass

rei nforcement

.

warp

(test direction)

Fill

(b)

Figure 20. G-10CR, warp test, 4K. SEM of fracture surface (37)

581
10

(0 ®

CL

S

O 6

X3

O ^

Compressive Strength

cr GU

4 K 76 K

40

co a

O

yj

Young’s Modulus

e

tu

IrVI

4 K

20 - 76 K

IlM'l 295 K

295 K

3

I-Hf

Warp

Fill

Tensile Strain to Failure

^TU

3 -

2 - 4 K

76 K

295 K

Figure 21. G-10CR. Compressive strength. Young's modulus,

tensile strain to failure vs. temperature (37).

582
o

*

LU

Bdo ‘smnaoiAi s.ONnoA

583

Figure

22.

G-10CR.

Young's

modulus

vs.

temperature

(37).
TEMPERATURE, K

Figure 23. G-10CR Compressive strength vs. temperature (37).

584
<0

h®

X.

3SV3HDNI %

585

Figure

24.

G-10CR.

Percent

increase

in

tensile

strain

to

failure,

Young's

modulus,

and

compressive

strength

at

cryogenic

temperatures

relative

to

room

temperature.

Calculated

from

results

in

reference

37.
Figure 25. G-10CR. Standard deviation vs. temperature (37).

586
Figure 26. G- 1 OCR . Coefficient of variation in percent vs. temperature (37).

587

STATIC AND DYNAMIC YOUNG'S MODULI OF TWO

FIBER-REINFORCED COMPOSITES

H. M. Ledbetter

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

Two uniaxially fiber-reinforced composites were studied: graphite/

epoxy and E-glass/epoxy. Dynamic Young's modulus was determined at kHz

frequencies using a Marx three-component oscillator. Comparison with

the static Young's modulus determined in compression shows agreement

within experimental error.

589

Elastic constants find many uses. In technological materials they

assume importance whenever problems arise concerning deflections, elastic

stability (buckling), excessive elastic deformation (jamming), thermoelastic

stress, residual stress, fracture, and so on. Ratios of elastic modulus

to density or to thermal conductivity provide well-known structural

design parameters. Elastic constants correlate empirically with many

practical properties such as strength, hardness, and wear. Because they

relate intimately to both the interatomic potential and the latticevibration spectrum, elastic constants relate in turn to many fundamental

solid-state phenomena such as diffusion, theoretical strength, and

lattice-defect energies. Elastic constants appear as coefficients in

thermodynamic equations of state where their relationships to free

energy, specific heat, and thermal expansivity become especially clear.

Since they can be measured accurately (typically within one percent or

5 less) and small changes can be detected easily (typically 1 in 10 or

better), elastic constants are useful physical properties for characteri zing

and comparing materials and for detecting changes in material properties.

Elastic constants are especially important for composites. These

materials exhibit elastic deformation at all stress levels up to rupture.

Methods for measuring elastic constants fall into two categories:

static and dynamic. Static measurements tend to be made in engineering

laboratories at high stress levels. Dynamic measurements tend to be

made in physics laboratories at very low stress levels. The dynamic

test frequency varies considerably : near 1 Hz for torsional -pendul um

and forced-vi bration measurements; up to several hundred kHz for resonance

measurements; up to 100 MHz for pulse-echo measurements.

591
Dynamic and static elastic-constant measurements remain unreconciled

among many engineers and scientists who use these properties. Engineers

tend to be more familiar with static measurements and tend to avoid

using results determined dynamically. This situation is unfortunate for

several reasons. First, dynamic results are more accurate by approximately

an order of magnitude. Second, dynamic measurements are quicker and

less expensive. Third, dynamic measurements on small specimens, including

sheet and wire, are relatively simple. Fourth, complete sets of elastic

constants, including anisotropy effects, are determined with little

di fficul ty.

Theoretically, dynamic-static elastic-constant differences arise

from three sources: the adiabatic-isothermal difference,^ the stressr 2 1 level (strain-level) difference, J and the role of higher-order elastic

F2l

constants at high-strain levels. J However, all these differences are

relatively small and relatively well understood so that the existing

static-dynamic controversy among practitioners surprises the theorist.

One feels the controversy will never be resolved except through

comparisons of carefully obtained static and dynamic results on the same

materials. Previously, the principal problem has been the dearth of

reliable static measurements.

This situation changed recently in round-robin Young 1 s-modul us

r 3 measurements made at six laboratories on two composites. J Table 1

shows the static Young's moduli, which are averages of measurements on

five specimens. Table 1 also shows the moduli determined dynamically at

60 and 30 kHz, respectively, by a three-component Marx-oscil lator method.

Three static methods were used: ASTM D695, ASTM D3410, and sandwich592
beam. The six laboratories participating in the round-robin testing are

designated by letters A through F in table 1. Two composites were

studied in this series of tests: (1) a unidirectional graphite/epoxy

(AS/3501) of high strength and medium modulus, and (2) a unidirectional

glass/epoxy laminate (E-glass/1002). Panels of both composites were

prepared by autoclave curing according to the manufacturer's recommendations

at the University of Delaware. The panels were cut into blocks that

were sent to participants who prepared specimens.

For both composites, results in table 1 show that the static

(determined compressively) and dynamic results agree within experimental

and/or material uncertainties. Thus, in these cases no basis exists for

distinguishing between static and dynamic Young's moduli. All three

theoretical factors listed above predict a dynamic elastic stiffness

that slightly exceeds the static. Although the graphite/epoxy results

show this relationship, the glass/epoxy results do not; there, the

seven-percent error for the static value masks this relationship.

Within our laboratory, we expect several further studies, theoretical

and experimental, on the static-dynamic elastic-constant question for

composites. This will include design and construction of apparatus for

accurate static values as a function of strain (stress).

Acknowl edgment

This study was supported partly by the DoE Office of Fusion Energy

and partly by the NBS Office of Standard Reference Data.

593
REFERENCES

[1] L. D. Landau and E. M. Lifshitz, Theory of Elasticity (Pergamon,

London, 1959), p. 19.

[2] F. D. Murnaghan, Amer. J. Math. 59_, 235 (1 937).

[3] N. R. Adsit, Compos. Tech. Rev. 2 (No. 2), 32 (1980).

[4] H. M. Ledbetter, in Nonmetal lie Materials and Composites at Low

Temperatures , (Plenum, New York, 1979), pp. 277-281.

594
Table 1 Static and Dynamic Young's Modulus,

in Units of GPa, for Two Fi ber-rei nforced Composites

AS/3501 [0°] E-glass/1002 [0°]

Static A 115 41

B 130 45

C 125 43

D 119 50

E 128 48

F 123 46

Static Avg. 123 ± 6 46 ± 3

Dynamic 130.4 ± 1 .3 44.3 ± 0.5

595

RADIATION EFFECTS ON COMPOSITES AND

NONMETALLIC MATERIALS

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

This discussion of radiation effects is part of a review paper on composites. The other parts of the review appeared last year, in Volume

III of this series.

597

The resistance of nonmetallic materials and composites to damage by

neutron and gamma irradiation of 4 K is an important consideration in

selecting insulating and structural support materials for superconducting

magnets of magnetic fusion energy systems. Current designs anticipate

that glass-reinforced polymers and other nonmetal lies could be exposed

to lifetime dose rates equivalent to about lO

1 *"* rads in some critical

areas.

1 The ability of conventional organic-matrix materials to withstand

such doses is significant concern.

A number of studies have been conducted on the effect of irradiation

on the room-temperature properties of nonmetallic insulators, but few

studies have extended into the cryogenic range. In addition to degrading

mechanical, thermal, and electrical properties, the low thermal conductivities

of the nonmetallics might also result in nuclear heating, which places

an added burden on the refrigeration system.

Extrapolating from room- temperature experience suggests that bisphenol A epoxy systems cured with aromatic amines should have better

radiation resistance than systems cured with dicyandiamide. Cycloaliphatic

epoxies may be superior to bishpenol A, polyimides should be better than

epoxies, and filled resin systems superior to unfilled resins, but confirmation is lacking. Complexities of the processes involved do not permit

a theoretical estimate to be made of the effect of irradiation at cryogenic

temperature on the various properties of nonmetallic insulators-- an

experimental data base must be acquired. Because the radiation resistance

of a polymer is strongly affected by the exact molecular structure, such

experiments must be carried on with materials characterized on the

2 molecular level

.

599
The Oak Ridge National Laboratory has recently screened the effect

of 4 K gamma and neutron irradiation on the stability of eight types of

organic insulation. The volume resistivities of a bisphenol A epoxy,

an inorganic-filled epoxy, an aromatic-polyimide sheet material, a

polyvinyl formal wire coating and an FR-5 type glass-epoxy laminate

O were found to be very little affected by a dose of 2 X 10 rads (y).

The effectiveness of a filled epoxy and of a B-staged glass-epoxy adhesive

for joining Cu-Be lap-shear tensile specimens was also relatively unaffected,

as was the flexural strength of a filled epoxy, a G-10 laminate, and an

FR-5 laminate.

When doses were increased to higher levels, the mechanical and

electrical properties of glass-reinforced epoxy laminates and filled

epoxy resins were found to be significantly deteriorated J Residual

flexural and compressive strengths of G-10CR and G-11CR laminates were

9 reduced to 10 to 15% of original values after exposure to 2.4 X 10 rads

(y) and to 6 to 10% after exposure to 1.0 X 10^ rads (y). These doses

20 2 20 were accompanied by fast neutron doses of 2.2 X 10 n/m and 8.7 X 10

2 n/m , respectively. The strength retention of the aromatic-ami ne-cured

G-11CR product appeared better than that of the dicyandiamide-cured

G-10CR after the lower dose, but no difference is noted after the higher

dose. A variant of the G-10CR product containing boron-free E-glass

reinforcement performed much like the G-11CR laminate.

The relatively high fraction of initial flexural modulus retained

by the laminates indicates that the glass reinforcement did not sustain major

damage. The low strain to failure indicates that deterioration of the matrix--

600
and possibly of the Fiber-matrix bond-accounts for most of the property

degradation.

The mechanical properties of the filled resins were relatively

unaffected by the lower level of radiation, but were substantially

degraded at the higher level. These materials appear to be essentially

at the end of their useful life after the higher level of radiation. 1

Despite the larger percentage of strength reduction, the high initial

strength of the glass/epoxy laminates permitted retention of sufficient

residual strength (100 MPa, 15 X 10

3

psi flexural strength; 55 MPa, 8 X

-3 10 psi compressive strength) for some magnet applications.

Except for one of the filled epoxies, no significant dimensional

changes were observed as a result of the irradiation.

Warming of the specimens to room temperature after 4-K irradiation

simulated periodic warming of the magnet structures throughout their

lifetime and incorporated additional damage anticipated from defect

recombination at higher temperatures--the Wigner energy.

It is apparent that much work remains to be done before reliable

radiation-resistant insulators are developed. Current programs are

examining the type of molecular damage incurred during irradiation of

bisphenol A epoxy-matrix laminates and testing is being extended to

include glass-reinforced laminates having other resin matrices.

601
REFERENCES

1. Coltman, R. R., Klabunde, C. E., Kernohan, R. H. and Long, C. J. "Radiation effects on organic insulators for superconducting magnets,"

Oak Ridge National Laboratory Report 0RNL/TM-7077 (1979).

2. Brown, B., "Low Temperature radiation effects on superconducting

magnets" submitted to Journal of Nuclear Materials (1980).

3. Kernohan, R. H., Coltman, R. R., and Long, C. J. "Radiation effects

on organic insulators for superconducting magnets," Oak Ridge National

Laboratory Report 0RNL/TM-6708 (1979).

602
NONMETALLIC MATERIALS

TECHNOLOGY TRANSFER

603

CONTRIBUTIONS TO TECHNOLOGY FORECAST '81 f

PLASTICS, CERAMICS AND COMPOSITES

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

^Submitted to Metal Progress

605

Nonmetallic composites perform critical structural , thermal, and

electrical insulation functions in large superconducting magnets required for magnetic fusion energy (MFE), magnetohydrodynamic (MHD)

,

rotating cryogenic machinery and energy storage systems required to meet

future energy requirements of the United States and other countries.

Large quantities of these material s--up to a thousand tons for a storage

magnet--must perform reliably over a thirty-year lifetime while subjected to large electromagnetic stresses, thermal stresses, and intense

radiation in some applications.

Studies at NBS have shown that the variability in cryogenic properties of conventional nonmetallic composites is a primary problem. This

variability reflects both the intrinsic embrittlement of the polymer

matrix at cryogenic temperatures and the extrinsic variability of cryogenic properties among laminates produced to similar commercial designations by a variety of producers.

NBS personnel are working with magnet designers, industrial laminators, and the polymer industry to establish standard grades of nonmetallic laminates for high technology cryogenic applications. The

principle is to establish specifications for components and manufacturing procedures that will minimize the property variability while providing a basis for systematic development of improved laminates in

response to demonstrated needs. This has resulted in the establishment of two commercial glass-fabric-reinforced epoxy laminate types

identified as NEMA G-10CR and G-11CR. The specifications have been

distributed to all U.S. laminators through the NEMA organization and

are presently being manufactured by three concerns.

607
Organic-matrix laminates are unable to retain required mechanical

and electrical properties after exposure to radiation levels on the

order of 10^° rads (y) plus a significantly fast neutron fluence at

4 K. Therefore, ceramics may be chosen over nonmetal lie composites for

insulation in MHD superconducting magnets, but ceramic use entails substantially higher magnet construction cost. On behalf of the Department

of Energy, NBS is cooperating with industry and with other federal laboratories to define the extent of the problem and to pursue development

of improved radiation resistant laminates. Studies at the Oak Ridge

National Laboratory have shown that the anticipated radiation levels

reduce the compressive and flexural strengths of conventional glassfabric epoxy laminates to 6 to 10% of their original values. Polyimide

matrices are presently being evaluated. Cycloaliphatic epoxy systems

and filled-resin laminates are primary candidates for future studies.

A significant detriment to the expanded use of composite materials in all areas of consumer technology is the difficulty of associating performance with the profusion of products identified by trade

names. Composite technology would greatly benefit from the establishment of an industry standards group who would establish a numerical

composite classification system based on the key elements affecting

performance. Such a system has been proposed by NBS, in which lower

bound mechanical, electrical and thermal properties are defined by a

number-1 etter combination defining type and flexibility of the matrix,

the type and configuration of the rei nforcement and the rei nforcement

volume fraction. The AISI designations for steel and the Aluminum

Association designations for aluminum alloys serve as models.

608
+COMMENT ON "MECHANICAL PROPERTIES OF PLASTIC COMPOSITES..

M. B. Kasen

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

+

To appear as a letter to the editor, Composi tes .

609

Comment on "Mechanical Properties of Plastic Composites Under Low Temperature Conditions" by K. Kaotani , Composi tes , V. 11, No. 2, April 1980,

pp. 87-94

The author's conclusion that composites fabricated with a highly

flexibilized epoxy matrix provide better cryogenic properties than those

fabricated with nonflexibil ized systems may be valid for laminates containing 32% by volume reinforcement, but is not valid for the higher

volume fractions typical of laminates in industrial usage. Literature

1 -3

reviews and work in our laboratory indicate that residual stresses

created by the large thermal contraction of highly flexibilized systems

negate any advantage provided by the intrinsically higher strain capability of the matrix when fiber contents approach 50 to 60% by volume

in glass, graphite, or aramid laminates. We find, for example, that

4

the flexible Resin 2 formulation decreases the 4-K uniaxial longitudinal tensile strength of an S- 901 glass-reinforced laminate to a level

about equal to that of a nonfl exi bi 1 ized matrix laminate at 295 K. The

transverse strength of a uniaxial graphite-reinforced laminate at 4 K

is reduced by about 60% by a highly flexibilized matrix, while the transverse failure strain is reduced about 90%. When used with aramid fiber,

the flexibilized matrix reduces the uniaxial transverse strength to an

extent rendering the laminate almost useless for structural applications.

Additionally, the intrinsic weakness of highly flexibilized resin

systems near room temperature has resulted in the catastrophic failure

of several structures designed for cryogenic use but necessarily subjected

to moderate stresses at room or slightly elevated temperatures.

611
The author quite correctly points out the limitation imposed on the

effective cryogenic strength of glass-reinforced laminates owing to the

strain limitations of conventional bisphenol A epoxies at such temperature (the "knee point" in the stress-strai n curve). But the solution

requires a resin system providing reasonably high room-temperature

strength and reasonably low thermal contraction as well as cryogenic

toughness. Possible approaches have been discussed in the recent literac ture. Until such systems are developed and proven, it appears most

prudent to utilize conventional resin systems for cryogenic laminates.

REFERENCES

1. Kasen, M. B., "Mechanical and thermal properties of filamentaryreinforced structural composites at cryogenic temperatures " Part I:

glass-rei nforced structural composites. Cryogenics 15, No. 6 (June

1975), pp. 327-349; Part II: advanced composites. Cryogenics , 15,

No. 12 (December 1975), pp. 701-719.

2. Kasen, M. B., "Cryogenic properties of fi lamentary-reinforced composites: An update," National Bureau of Standards, Boulder, Colorado

(in preparation )

.

3. Schramm, R. E. and Kasen, M. B., "Cryogenic mechanical properties

of boron-, graphite-, and glass-reinforced composites," Material

s

Science and Engineering , 30 (1977), pp. 197-204.

4. Softer, L. N. and Molho, R., "Cryogenic resins for glass-filamentwound composites," NASA CR-72114 (Final).

5. Nonmetal lie Materials and Composites at Low Temperatures , A. F.

Clark, R. P. Reed, and G. Hartwig, eds. (Plenum Press, 1979).

612
THE MANUFACTURE OF RADIATION RESISTANT LAMINATES

J. R. BENZINGER

CORPORATE TECHNOLOGY

SPAULDING FIBRE COMPANY, INC.

TONAWANDA, NEW YORK 14150

PRESENTED AT THE NBS/DOE WORKSHOP ON MATERIALS AT

LOW TEMPERATURE, OCTOBER 8, 1980, VAIL, COLORADO

613

THE MANUFACTURE OF RADIATION RESISTANT LAMINATES

J. R. BENZINGER

Corporate Technology

Spaulding Fibre Company, Inc.

Tonawanda, New York 14150

INTRODUCTION

There have been many studies on the irradiation effects of

polymers and insulators, but only a few on high pressure thermosetting laminates'^. There was virtually no published information

on radiation at cryogenic temperatures on polymers, let alone

laminates, until the recent work at Oak Ridge on the Spauldite©

CR grade laminates, some particle filled epoxies and films of

2 Mylar and Kapton . One conclusion of the 1979 ORNL report wr as

in

tnat , at a dose ot 1 x 10*“ rads, particle filled epoxies are at

"end of life" in terms of mechanical strength while CR grade

laminates, although showing severe degradation, still retained

usable strength. Further testing was suggested to determine

margins of safety. In some locations of the magnet, such as

adjacent to the injector of neutron beams, higher doses would

require additional shielding or redesign at considerable

expense . A more prudent approach it is claimed would be to

improve the radiation resistance of the insulation.

Inorganic insulators like ceramics are much more resistant

to radiation but tend to be very brittle. This rules them out

615
for most fusion reactor applications since these magnets are

subject to extreme mechanical and magnetically induced forces^.

Consequently, the critical applications can only be satisfied

by the use of tough, high pressure thermosetting laminates^.

It is the purpose of this paper to demonstrate the capability

of industry to modify high pressure thermosetting laminates for

improved radiation resistance.

Mechanical strength testing of some of the experimental

G-10CR variants designed for improved radiation resistance

yielded some surprising results. Also to be reviewed are the

design rationale, materials of construction, and method of

manufacture of some polyimide and epoxy laminates prepared for

MIT and ORNL irradiation effect studies. In conclusion, cost

comparisons, availability and anticipated manufacturing problems

associated with some raw materials and laminates will be discussed .

MIT AND ORNL PROGRAMS

In response to requests from H. Becker and E. Erez of MIT

and J. Long of ORNL, several experimental laminates were prepared

in the laboratory with the intent to improve radiation resistance.

Some of them were variants in some manner to Spauldite© G-10CR

or G-11CR.

The program at ORNL, sponsored by DOE is an ongoing special

purpose materials program on "Radiation Effects On Organic

616
Insulators For Superconducting Magnets". The MIT work was

performed in the MIT Plasma Fusion Center also under DOE auspices. This work was part of an investigation of material

properties for an ignition test reactor (ITR) subject to high

neutron fluence and gamma radiation^. The radiation effect

results on the laminates of these programs will be reported

by their respective authors at the NBS/DOE Workshop On

Materials At Low Temperature at Vail, October 7-9, 1980.

MATERIALS OF CONSTRUCTION

Fabric Reinforcements

Listed in Table I is a description and cost of the

woven fabrics used as reinforcements in the experimental

nates. The low cost plain weave E glass 7628 A-1100 was

as a control based on an extensive test data base and a

7 history of successful use in laminates .

five

lamiselec ted

long

S glass was selected not only because of its high strength,

but also because it is boron free. Table II lists the chemical

composition of the various glasses. E glass which contains

boron may present a problem in reactors in that it retains

residual radioactivity for longer periods of time than boron

free glasses. Quartz was chosen because it tested the best of

several inorganic insulators for resistance to irradiation.

The standard finish available on the commercial S and

quartz fabrics was A-1100, the same finish as used in the CR

617
grade laminates. At least one investigator has concluded that

A-1100 was the best of a variety of coupling agents evaluated .

It is suggested that some future program include an update

study of coupling agents when time permits ordering and preparing

special samples.

Although all of the five fabrics used in this program were

commercial grades, availability and cost of quartz and S glasses

could present problems in delivery. As for manufacturing, there

are less problems associated with the manufacture of high

pressure laminates with plain weaves. However, the satin weaves

tend to perform better in low pressure molding where curvatures

are involved. See Figure 1 for a schematic of the comparisons

of construction.

Resins and Curing Agents

It has been reported by several early researchers that

aromatic compounds such as polyimides and aromatic epoxies have

g good radiation resistance . They are many times more stable

than aliphatic compounds to radiation owing primarily to the

greater bond strength and the resonance energy of the aromatic

structure. Substituted aromatics are more resistant than nonsubstituted aromatics and saturated aliphatics are more stable

than unsaturated aliphatics. Other investigators have reported

that the aromatic amine curing agent diaminodipheny 1 sulphone

was best when used with epoxies of the diglycidyl ethers of

bisphenol A type an^ that diaminodiphenyl methane gave the best

618
results when used with multi-functional glycidyl amines 10

. It

was also observed that the more highly functional and crosslinked polymers that perform well at high temperatures also tend

to fair best at high levels of radiation. Listed in Table III

are the epoxies, curing agents and polyimides selected for the

program.

The polyimide, Kerimid 601, was selected for its ease of

processing, modest cost as far as polyimides are concerned and

because it has been reported to be unaltered after a 10 10 rad

exposure. The other polyimide used in the program was DuPont

NR-150B2 polyimide precursor solution. Shortly after the

experiments were underway, DuPont withdrew this resin from the

market

.

LAMINATE PREPARATION

The laminates prepared for the MIT and ORNL programs are

listed in Tables IV and V. Most of the laminates were very

thin and consisted of only two or three sheets of prepreg.

For this reason the laminate specimens were prepared by laboratory "hand treating" rather than using the pilot facility.

This process involved hand dipping of the reinforcements

in solutions consisting of the resin polymer, its curing agent

and compatible solvents. The resin treated webs were then

passed through motor driven squeeze rolls. The rolls were

preset so that pressure was applied to the treated web forcing

the resin into the interstices of the fabric. The force of the

G19
rolls also squeezed off excess resin, leaving a uniform application .

After saturation, the treated webs were placed in an aircirculating oven at 280-300° F. The purpose of the drying

operation for the polyimides was to remove the solvent. In the

case of epoxies and phenolics, additional polymerization takes

place in the drying oven in addition to the removal of solvent.

All of the prepregs were subsequently pressed between steel

plates at a pressure of 500 psi and a temperature of 360°F with

the exception of one of the Kerimid laminates which was pressed

at 4 40 ° F

Two of the polyimide laminates supplied to ORNL were post

cured at 450°F for five hours to insure adequate cure. The

epoxy based laminates were not post cured. Results of some

radiation effects by R. Sheldon on one wet lay-up epoxy laminate

indicated that higher temperature post cure tended to improve

8 radiation resistance . Subsequent designed experiments at

different post cure temperatures by D. Evans indicated no signi9 ficant differences could be detected in radiation resistance . Therefore, for this reason and because the process is irreversible, the epoxy laminates were not post cured. Based on

initial radiation results, a post cure study, if deemed advisable,

could be undertaken simply by post curing replicate samples of

the laminates.

Also included for controls were the commercially available

laminate, Spauldite® G-11CR and a phenolic cured epoxy novolac.

620
G-10CR VARIANTS

GEM-CR, Quartz Filled

As part of an ongoing research program at Spaulding Fibre

Company, Inc. several variants of G-10CR and G-11CR were investigated with a view toward improvements in radiation resistance,

mechanical strength and cost. One of these was a continuous

filament E glass epoxy mat laminate (GEM-CR) which was constructed

using the G-10CR resin system.

Last year Dr. Olinger of OCF reported on the good results

he had achieved at liquid nitrogen temperatures using mat epoxy

laminates'^ . Also several investigators have reported improved

radiation stability by incorporating particle fillers in epoxy

1 2 resin castings and moldings . Other improved properties

reported for filled resins are less thermal contraction, reduced

cracking or crazing, increased or reduced thermal conductivity

1

based on the type and amount of filler and cost . Therefore,

a combination of reinforcement, resin and filler seemed like a

good approach to improve radiation resistance and other properties of laminates as was the case with castings.

The initial experiment was conducted with glass mat rather

than fabric for three reasons: (a) lower cost, (b) random and

wide spacing of the fibers for reception of the filler, and (c)

prior success in GPO laminates. Although several fillers were

considered as possible candidates, see Figure 2, the filler

selected for experimentation was Quartz (SiC^) because, as

621
stated previously, it tested the best of several inorganic

insulators for resistance to irradiation 14

. Shown in Figure V

are some mechanical strength properties of both filled and

unfilled GEM-CR. The properties of the filled GEM-CR laminate

were not as good as the unfilled GEM-CR, but this is believed

due to the high loading of SiC^, 80 parts to 100 parts of resin.

A high loading was used to reduce costs as well as improve

radiation stability. Dr. Kasen has suggested that while the

optimum amount is unknown, it is possible that a small amount,

well distributed, could be effective.

The bar graph in Figure 3 compares the compressive strength

edgewise at ambient (293 K) and liquid nitrogen (77 K) temperatures of the unfilled GEM-CR laminate versus a NEMA GPO-1.

(glass mat polyester). The GEM-CR variant like G-10CR increases

considerably in strength at cryogenic temperatures. The bar

graph in Figure 4 is similar to Figure 3 in all respects except

that it compares interlaminar shear (unsupported) as a function

of temperature.

S Glass And High Glass Content

Some mechanical properties of the S glass variant of G-10CR

are shown in Table V compared with the average test data of

Spauldite® G-10CR. This laminate was made identical to G-10CR

with the only variant being S in place of E glass. The weave,

finish and resin were the same as G-10CR. The two properties

that stand out are compressive strength normal to laminations

G22
and bond strength. Both tests when repeated gave the same

identical results. The high E glass content variant of G-10CR

demonstrates another way to achieve higher strengths.

Both of the above laminates, one because of the absence of

boron oxide and other because of less organic content, would be

expected to provide improved radiation resistance.

Polyimide And Phenolic-Epoxy

Also listed in Table V are the properties of the polyimide

laminate submitted to Joe Long at Oak Ridge for irradiation

effects. This laminate is grouped with an aniline phenolic

cured epoxy novolac epoxy. Both laminates contain high temperature, highly aromatic polymers which like G-11CR are considered

key elements for resistance to radiation. These resins are

harder and tend to be a little brittle as evidenced by the

reduced interlaminar shear and bond strengths when compared with

G-10CR (epoxy resin) laminates.

G-10CR Quartz Filled

Based on the good results achieved with the particle filled

glass epoxy mat laminate, a quartz filled variant of G-10CR was

planned. Several changes were made, one of which was to reduce

the amount of quartz (SiC^) added to the epoxy resin from 80 to

40 parts by weight to 100 parts epoxy. The selection of a finer

particle size SiC^ filler, half the size used in the mat, was

considered another improvement. In order to improve the bond of

the filler to the epoxy with resultant expected increase in

523
strength a filler was obtained that had been pretreated with

A-1100 coupling agent.

Another reason to explain the lower strength of the filled

GEM-CR laminate is that the percentage of glass fiber mat reinforcement was reduced by the exact amount of the loading of

filler while the amount of resin remained the same as in the

unfilled laminate. In the case of the SiQ^, filled fabric

laminate, the resin, rather than the reinforcement was displaced

by the filler, and the amount of glass fabric reinforcement was

not reduced but kept the same as the unfilled G-10CR.

Test data on mechanical strength properties of this latest

variant of G-10CR is not available for inclusion in this report,

but will be reported verbally at the workshop. It should be

emphasized that all of the laminates reviewed in this paper,

with the exception of G-10CR and G-11CR, are laboratory developments and not commercially available at this time. Also, most

of the reported data are based on one time testing of laboratory

specimens

.

CONCLUSION

It has been demonstrated that laminates can be modified in

many ways to make them more resistant to radiation. The success

of these endeavors will be largely unknown until and when the

laminates can be properly tested and evaluated. Higher temperature irradiations can serve as a guide, but as concluded by B.

Brown only irradiations performed near the temperature of the

624
magnet are realistic'*"'’. If it is deemed necessary that a more

comprehensive program be established, it seems more than likely

that a significant improvement in radiation resistance can be

successfully achieved with high pressure laminates.

ACKNOWLEDGEMENT

The author wishes to gratefully thank E. Erez of MIT for

the test data used to generate the graphs in Figures 3 and 4.

REFERENCES

1. M. G. Young, ’’Radiation Effects On Laminated Plastic

Materials’’, UD-NEMA Facility, April 1965

2. R. H. Kernohan, R. R. Coltman and C. J. Long, ’’Radiation

Effects On Organic Insulators For Superconducting Magnets”,

Oak Ridge National Laboratory Report 0RNL/TM-6705 (1979)

3. R. T. Santaro. J. S. Tang. R, G. Alsmiller, Ji . and J. M.

Barnes, Nucl . Tech. _37 , 65 (1978)

4. R. P. Reed, F. R. Fickett, M. B. Kasen and H. I. McHenry,

’’Materials Studies For Magnetic Fusion Energy Applications

At Low Temperature - 1”, Report No. NBSIR 78-884, National

Bureau of Standards, Boulder, Colorado (1978), p. 243

5. M. B. Kasen, G. R. MacDonald, D. H. Beekman and R. E.

Schramm, "Mechanical, Electrical And Thermal Characterization

Of G-10CR And G-11CR Glass Cloth/Epoxy Laminates Between

Room Temperature And 4 K”, Advances In Cryogenic Engineering,

V 26 Plenum Press N.Y. (1980)

625
6 . E. A. Erez, H. Becker, "Radiation Damage In Thin Sheet

Fiberglass Insulators", Non-metallic Materials And Composites At Low Temperature - II, Plenum Press 1981

7. J. R. Benzinger, "Advances In Cryogenic Engineering",

Volume 26 Plenum Press, New York 1980 pp . 252-258

8. R. Sheldon and G. B. Stapleton, "The Effect Of High Energy

Radiat ion On The Mechanical Properties Of Epoxy Resin

Systems Used For Particle Accelerator Construction",

Rutherford Laboratory, Chilton, Didcot, England (1968)

9. M. H. Van de Voorde, IEEE Trans. Nucl. Sci. 18 , 784 (1970),

and IBID 2j0, 693 (1973)

10. D. Evans, J. T. Morgan, R. Sheldon, G. B. Stapleton, "PostIrradiation Mechanical Properties Of Epoxy/Glass Composites"

RHEL/R200 Chilton, Berkshire, England, 1970

11. J. Olinger, "Cryogenic Grade Glass Mat Epoxy Laminates”,

NBS/DOE Workshop On Materials At Low Temperature", Vail,

Colorado, October 1979

12. H. Brechna, "Effect Of Nuclear Radiation On Organic

Materials; Specifically Magnet Insulations In High-Energy

Accelerators", SLAC Report No. 40, 1965

13. G. Claudet, F. Disdier and M. Locatelli, "Interesting Low

Temperature Thermal And Mechanical Properties Of A

Particular Powder-Filled Polyimide", Non-Metallic Materials

And Composites At Low Temperatures, Plenum Press pp. 131-

140, 1979

626
14 . C. L. Hanks and D. J. Hamman, "Report On The Effect Of

Radiation On Electrical Insulating Materials", REIC Report

No. 46, Battelle Institute, Columbus, Ohio (1969)

15. B. S. Brown, "Radiation Effects In Superconducting Fusion

Magnet Materials", May 1980, submitted to Journal Of Nuclear

Materials

LIST OF TABLES

TABLE I Description Of The Woven Fabric Reinforcements

TABLE II Chemical Composition Of The Reinforcements

TABLE III Polymeric Materials Of Construction

TABLE IV Fabric Reinforced Polyimide Laminates Prepared

For Irradiation

TABLE V Fabric Reinforced Epoxies Prepared For

Irradiation

TABLE VI Mechanical Strength And Physical Properties Of

Some G-10CR Variants

TABLE OF FIGURES

Figure 1. Comparison Of Weaves, Style 7628 And Style 181

Figure 2. Relative Radiation Resistance Of Inorganic

Insulating Materials, Based Upon Changes In

Physical Properties

Figure 3. Compressive Strength Edgewise As A Function Of

Temperature Spauldite® GEM-CR Versus NEMA GPO-1

527
Figure 4. Interlaminar Shear Strength ASTM

(Unsupported) As A Function Of

GEM-CR Versus GPO-1

Guillotine

Temperature

628
TABLE

I

-

DESCRIPTION

OF

THE

WOVEN

FABRIC

REINFORCEMENTS

G29

Control

fabric

used

in

Spauldite®

G-11CR
TABLE II - CHEMICAL COMPOSITION OF THE REINFORCEMENTS

COMPONENTS, % BY WEIGHT E GLASS S GLASS QUARTZ

Silicon dioxide 52-56 64-66 99.9

Calcium oxide 16-25 —

Aluminum oxide 12-16 24-26

Boron oxide 8-13 —

Sodium ^ potassium oxide 0-1 —

Magnesium oxide 0-6 9-11

Density, g/cc 2.54 2.48 2.65

G30
TABLE

III

-

POLYMERIC

MATERIALS

OF

CONSTRUCTION

x

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631
TABLE

IV

-

FABRIC

REINFORCED

POLYIMIDE

LAMINATES

PREPARED

FOR

IRRADIATION

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c^j £

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TABLE

V

-

FABRIC

REINFORCED

EPOXIES

PREPARED

FOR

IRRADIATION

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634

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TECHNOLOGY TRANSFER

639

NBS-DoE WORKSHOP: MATERIALS AT LOW TEMPERATURES

October 7-9, 1980, Vail, Colorado

N. J. Simon

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado

ABSTRACT

The subject matter and format of this annual workshop are discussed.

During fiscal year 1980, this workshop constituted the chief formal activity of the technology transfer program. However, starting with the

present fiscal year, a second major activity will be undertaken: the

production of handbook pages for low-temperature mechanical and physical

properties of materials for magnets in fusion energy systems.

641

The workshop discussions concerned structural and insulating materials to be used at low temperatures in superconducting magnets for

fusion reactors and MHD systems. Each year, a major purpose of the

workshop is to bring together representatives from the community of

users and suppliers of materials for low temperature magnet applications

with those engaged in development and characterization of these materials.

Participants this year came from industry, the national laboratories,

the magnetic fusion energy programs, NBS, DoE and several universities.

A list of attendees is given in the following pages. The technical presentations were informal, with many opportunities for comments and questions.

Initially, overviews of the large superconducting magnet programs

for fusion energy and MHD were given. Presentations regarding the Large

Coil Projects and the ORNL test facility followed. Then, recent measurements

of properties of structural alloys, welds and castings were presented.

In the next session, fracture, fatigue and radiation damage in composites

were considered. In the last session, codes and standards for fusion

energy magnets were discussed. A detailed list of all the technical

papers presented is attached. The proceedings of the workshop are not

published, but abstracts or short summaries of all the presentations are

provided to those who attend.

Meetings of several advisory groups were held during the workshop.

For example, the CCP conductor sheath advisory committee met, as did the

joint NBS-industry group which is working on specifications for CR-grade

laminates.

643
PROGRAM

NBS/DoE Workshop on Materials at Low Temperatures

October 7-9, 1980

Vail, Colorado

Tuesday, October 7

Session 1 (9:00-12:00): Large Superconducting Magnet Programs

R. P. Reed, Chairman

DoE Office of Fusion Energy - Overview

NBS Program

Magnetic Fusion Test Facilities, A and B Engineering Test Facility

MHD Magnets

D. Beard, DoE

R. P. Reed, NBS

E. N. C. Dalder, LLL

R. J. Hooper, 0RNL

V. Der, DoE

Session 2 (2:00-5:00): Large Coil Projects

D. Beard, Chairman

Oak Ridge National Laboratory

General Dynamics

General Electric

Westi nghouse

Japan Atomic Energy Research Institute

Kernforschungszentrum Karlsruhe

C. J. Long, 0RNL

J. L. Christian, GD

C. L. Linkinhoker, GE

C. J. Heyne, W

E. Tada, JAERT

A. Ulbricht, KfK

Wednesday, October 8

Session 3 (8:30-12:00): Structural Alloys, Welds and Castings

H. I. McHenry, Chairman

Cr-Ni-Mn-N Stainless Steels

lanon^rA C ^ m W ^ Jw bUU I CO

Conductor Sheath Materials

Welding Research at Lawrence Berkeley

Stainless Steel Weldments

Aluminum Weldments

Stainless Steel Castings

R. L. Tobler, NBS

R. Ogawa, KOBE

R. Gold, W

J. W. Morris, LBL

I. A. Whipple, NBS

J. W. Elmer, NBS

R. Finch, £SC0

Session 4 (1:30-5:00): Insulating and Structural Composites

M. B. Kasen, Chairman

Fracture Modes of G-10CR in Compression

Fatigue of Fibrous Glass/Epoxy Laminates in LN^

Multiaxial Fatigue Damage in Fiberglass

Radiation Damage in Fiberglass Insulators

Radiation Effects in Organics at 4 K

Manufacturing of Radiation Resistant Laminates

Low Temperature Irradiation Studies

B. A. Beck, NBS

J. L. Olinger, Owens Corning

S. Wang, UI

H. Becker, MIT

C. J. Long, 0RNL

J. Benzinger, Spaulding

B. Brown, ANL

6:00 - 8:00 P.M. COCKTAIL PARTY

Thursday, October 9

Session 5 (8:30-12:00): Codes and Standards

E. N. C. Dalder, Chairman

Structural Design Basis for Superconducting Magnets

The Development of Codes and Standards

Cooperative Research for Codes and Standards

Code Approaches to Fracture Control

Material Suppliers Viewpoint

Filler Metal Suppliers Viewpoint

H. Becker, MIT

W. E. Cooper, TES

M. Prager, MPC

H. I. McHenry, NBS

E. H. Espy, Armco

D. J. Kotecki, McKay

644
NBS/DOE

WORKSHOP

ATTENDANCE

October

6-9,

1980

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645
646
ORGANIZATIONAL CONTACTS

The people responsible for the various major aspects of the program

are listed below so that specific technical questions may be directed to

them.

Department of Energy, Office of Fusion Energy, Washington,

Program Monitor D. Beard

D.C. 20545

(301) 353-4958

(FTS) 233-4958

National Bureau of Standards, Boulder, Colorado

Program Manager R. P. Reed

80303

(303) 497-3870

(FTS) 320-3870

Welding and Casting H. I. McHenry

T. A. Whipple

E. L. Brown

(303) 497-3268

497-5795

497-5790

Elastic Properties

Nonmetal 1 ics

Structural Alloys

Technology Transfer

H. M. Ledbetter

M. B. Kasen

R. P. Reed, R. L. Tobler

N. J. Simon

(303) 497-3443

(303) 497-3558

(303) 497-3421

(303) 497-3687

ESC0 Corporation, 2141 N.W. 25th

Program Manager

Avenue, P. 0. Box 10123, Portland, Oregon 97210

L. Finch (503)228-21 41 ,X520

University of Wisconsin, Engineering Research Building

Program Manager D. Larbalestier

, Madison, Wisconsin

(608)262-2194

53706

647
ACKNOWLEDGMENT

The editors wish to express their appreciation to JoAnne Wilken for

preparing much of the manuscript. They are also grateful for the assistance of Carole Montgomery and Vickie Grove.

Marilyn Stieg removed many of the inconsistencies, grammatical

errors, and awkward usages of language. The editors are responsible

for those errors that remain.

648
NBS-114A (REV. 2-8C) U.S. DEPT. OF COMM.

4.

1 . BIBLIOGRAPHIC DATA

SHEET (See, nstructions)

PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No. 3.

NBSIR 81-1645

Publ ication Date

April 1981

TITLE AND SUBTITLE

Materials Studies for Magnetic Fusion Energy Applications at Low Temperatures- IV

5. AUTHOR(S)

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

6. PERFORMING ORGANIZATION (If joint or other than NBS. see in struction s) 7. Contract/Grant No. NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City. State. ZIP)

Department of Energy

Office of Fusion Energy

Washington, D.C. 20545

10.

SUPPLEMENTARY NOTES

Related earlier reports: NBSIR 80-1627, NBSIR 79-1609, NBSIR 78-884

Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11.

ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here)

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

that will facilitate design and development of cryogenic structures for the superconducting magnets of magnetic fusion energy power plants and prototypes. The program

was conceived and developed jointly by the staffs of the National Bureau of Standards

and the Office of Fusion Energy of the Department of Energy; it is managed by NBS and

sponsored by DoE. Research is conducted at NBS and at various other laboratories through

subcontracts with NBS. The reports presented here summarize the fourth year of work on the low temperature materials research program. Highlights of the results are presented

first. Research results are given for the four main program areas: structural alloys,

weldings and castings, nonmetal 1 i cs , and technology transfer. Objectives, approaches,

and achievements are summarized in an introduction to each program area. The major

portion of the program has been the evaluation of the low temperature mechanical and

physical properties of stainless steel base 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 and standardization of

nonmetal 1 ics , primarily industrial laminates, for low temperature applications and on the

measurement of their properties at cryogenic temperatures. A brief description is given

of the fourth NBS/DoE Vail workshop held in October 1980.

12.

KEY WORDS (Six to twelve entries ; alphabetical order; capitalize only proper names ; and separate key words by semicolons)

Alloys; aluminum alloys; castings; composites; cryogenic properties; low temperature;

mechanical properties; nonmetallics; physical properties; stainless steels;

structural i ^ t^i vi a—art alloys: a , a y j as ^ welding. " v- » >» i . ^

* 13. AVAILABILITY 14. NO. OF

PRINTED PAGES

Uni imited

Q31 For Official Distribution. Do Not Release to NTIS

Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 604

20402. 15. Price

QX] Order From National Technical Information Service (NTIS), Springfield, VA. 22161 $41.00

USCOMM-DC 6043-P80 h U.S. GOVERNMENT PRINTING OFFICE:1981—780-205 / 386