DP-1643

D i s t r i b u t i o n Category: UC-25

HYDROGEN COMPATIBILITY HANDBOOK ^^-i^^s

FOR STAINLESS STEELS DE83 017051

GEORGE R. CASKEY. JR.

Approved by

R. L. Folger, Research Manager

Hydrogen and Ceramic Technology Division

Publication Date: June 1983

DISCLAIMER

This report was prepared as an account of work sponsored by an agency of the United States
Government. Neither the United States Government nor any agency thereof, nor any of their
employees, makes any warranty, express or implied, or assumes any legal liability or responsi-
bility for the accuracy, completeness, or usefulness of any information, apparatus, product, or
process disclosed, or represents that its use would not infringe privately owned rights. Refer-
ence herein to any specific commercial product, process, or service by trade name, trademark,
manufacturer, or otherwise does not necessarily constitute or imply its endorsement, recom-
mendation, or favoring by the United States Government or any agency thereof. The views
and opinions of authors expressed herein do not necessarily state or reflect those of the
United States Government or any agency thereof.

E. I. du Pont de Nemours & Co.

Savannah River Laboratory
Aiken, SC 2 9 8 0 8

DISTRIBUTION Of THIS COCUMENl IS UHLIfMEO

 DISCLAIMER

This report was prepared as an account of work sponsored by an

agency of the United States Government. Neither the United States
Government nor any agency Thereof, nor any of their employees,
makes any warranty, express or implied, or assumes any legal
liability or responsibility for the accuracy, completeness, or
usefulness of any information, apparatus, product, or process
disclosed, or represents that its use would not infringe privately
owned rights. Reference herein to any specific commercial product,
process, or service by trade name, trademark, manufacturer, or
otherwise does not necessarily constitute or imply its endorsement,
recommendation, or favoring by the United States Government or any
agency thereof. The views and opinions of authors expressed herein
do not necessarily state or reflect those of the United States
Government or any agency thereof.
 DISCLAIMER

Portions of this document may be illegible in

electronic image products. Images are produced
from the best available original document.
I
ABSTRACT

This handbook compiles data on the effects of hydrogen on the

mechanical properties of stainless steels and discusses this data
within the context of current understanding of hydrogen compati-
bility of metals. All of the tabulated data derives from continu-
ing studies of hydrogen effects on materials that have been con-
ducted at the Savannah River Laboratory over the past fifteen
years. Supplementary data from other sources are included in the
discussion. Austenitic, ferritic, martensitic, and precipitation
hardenable stainless steels have been studied. Damage caused by
helium generated from decay of tritium is a distinctive effect that
occurs in addition to the hydrogen damage from tritium which is the
same as for the other hydrogen isotopes protium and deuterium. The
handbook defines the scope of our current knowledge of hydrogen
effects in stainless steels and serves as a guide to selection of
stainless steels for service in hydrogen.

-3-f
 A/
'''••h

/o
••^"'t4f

I
CONTENTS

INTRODUCTION 11

SUMMARY 13

HYDROGEN DAMAGE 16

Mechanical Behavior 17

Material and Processing Variables Affecting Hydrogen Damage 18

Composition 18
Heat Treatment 23
Mechanical Processing 26
Effect of Martensite 35

Environmental Variables Affecting Hydrogen Damage 37

Hydrogen Pressure 37
Test Temperature 37
Stress State and Strain Rate 41
Notch Tensile Tests 46 '
J-Integral Analysis 46
Static Loading 50
Strain Rate 53

Fractography 53

Fracture Modes 53

Microvoid Coalescence 54
Twin-Boundary Parting 54
Transgranular Cleavage 56
Intergranular Separation 56
Interphase Separation 63

Correlation of Fracture Mode with Composition 65

Helium Embrittlement 67
- 5-
CONTENTS, Contd

CONCLUSIONS 75

ACKNOWLEDGEMENTS 77

ALLOY DATA SHEETS 78

Alloy Index 78

Iron-Chromium-Nickel Alloys 78

Iron-Chromium-Nickel-Manganese Alloys 78

Precipitation Hardenable Alloys 78

High Purity Alloys 78

REFERENCES 124

APPENDIX A. DEFINITIONS 131

APPENDIX B. CONVERSION TABLES 133

Temperature Conversions 133

English/Metric (Sl) Stress Conversion Factors 134

English/Metric (Sl) Fracture Toughness Conversion Factors 1

English/Metric (SI) Impact Energy Conversion Factors 137

APPENDIX C. MECHANICAL TEST SPECIMENS 138

Smooth Bar Tensile Specimens 138

Notched Bar Tensile Specimens 140

Tensile Tube Specimen 142

Single Edge Notched Specimen 143

C-Shaped Fracture Mechanics Specimen 145

Impact Specimen 147

APPENDIX D. HEAT ANALYSES 148

LIST OF TABLES

1 Ranking of Hydrogen Damage 14

2 Hydrogen Damage Susceptibility 14

3 Mechanical Properties of HERF and Annealed Alloys in

High-Pressure Gas 31

4 Effect of Martensite on Subsequent Environmental Hydrogen

Damage in Type 304L Stainless Steel 36

5 Temperatures for Strain-Induced Martensite Formation 43

6 Effect of Hydrogen Charging on Tensile Properties of

Type 304L Stainless Steel 49

7 Ranking of Hydrogen Compatibility (HERF Stainless

Steels) 49

8 Stress Necessary for Slow Crack Growth in Type 304L

Stainless Steel 52

9 Occurrence of Microcracks in Hydrogen-Saturated Austenitic

Stainless Steels 58

10 Mechanical Properties of HERF Type :i04L Stainless Steel 70

11 Mechanical Properties of HERF Nitronic® 40 Stainless Steel 70

12 Tritium and Helium Effects on Fracture Toughness

of Stainless Steels 73

- 7-
LIST OF FIGURES

1 Environmental Hydrogen Damage in Fe-Cr-Ni Alloys at Room

Temperature and 69 MPa Hydrogen 19

2 Stacking Fault Energies in Fe-Cr-Ni Alloys at Room

Temperature 21

3 Environmental Hydrogen Damage in Fe-Cr-Ni-Mn Alloys at Room

Temperature and 69 MPa Hydrogen 22

4 Hall-Petch Plot for Tensile-Test Specimens of Type 304L

Stainless Steel at a Plastic Strain e = 0.05 24

5 Hall-Petch Plot of Yield Strength for Type 304L Stain-

less Steel Tested in High-Pressure Helium or Hydrogen
Environments 25

6 Effect of Prestrain on Tensile Strength of Type 304L Stainless

Steel Tubes in Air or Hydrogen 27

7 Effect of Prestrain on Elongation of Type 304L Stainless Steel

in Air or Hydrogen 28

8 Effect of Prior Cold Work on Hydrogen Damage Susceptibility of

Stainless Steels 30

9 Effect of Processing on Hydrogen Damage in Nitronic® 40

Elongation 32

10 Effect of Processing on Hydrogen Damage in Nitronic®

40 Strength 33

11 Orientation of V-Notch with Respect to Forging Flow Lines

in Tensile-Test Specimens Machined from Bars of HERF
Stainless Steels 34

12 Effect of Hydrogen Pressure on Ductility of Type 304L

Stainless Steel Tested to Fracture in Hydrogen at Room
Temperature 38

13 Dutility Minima in Fe-Cr-Ni Alloys Charged with

Deuterium at 69 MPa and 620 K for Three Weeks 39

14 Ductility Minima in Fe-Cr-Mn-Mi Alloys Charged

with Deuterium at 69 MPa and 620 K for Three Weeks 40
LIST OF FIGURES, CONTD

15. Isoductility Diagram for Hydrogen-Charged Fe-Cr-Ni Alloys 42

16 Effect of Hydrogen on Strain-Induced Martensite Formation in

Type 304L Stainless Steel 44

17 Effect of Hydrogen on Strain-Induced Martensite Formation in

Tenelon® and Nitronic® 40 45

18 Relative Effect of Hydrogen on Properties of Sensitized

Type 304L Stainless Steel 47

19 Environmental Hydrogen Damage in Nitronic® 40 Stainless Steel

at Room Temperature 48

20 Hydrogen Effects on Crack Growth in Stainless Steels 51

21 Dimpled Fracture of Type 304L Stainless Steel 55

22 Twin-Boundary Parting in type 304L Stainless Steel. High

Energy Rate Forged. Tested at 2(K) K 55

23 Microcracks Along Boundaries of Annealing Twins 57

24 Variation of Facet Appearance with Test Temperature in

Type 304L Stainless Steel 59

25 Multiple Crack Nucleation Along Boundaries 60

26 Transgranular Cleavage in 17-4 PH Precipitation-

Hardenable Stainless Steel 61

27 Intergranular Separation in Nickel and Inconel® 718 61

28 HERF Nitronic® 40, Orientation 1. Intergranular Fracture in

Hydrogen 62

29 Fracture Along Interphase Interfaces in

Austenitic Steels 64

30 Composition Regimes for Fracture Modes Observed in HAF of

Iron-Chromium-Nickel Alloys 66

31 Helium bubbles in Type 304L Stainless Steel Tritium Charged,

Aged, and Annealed at 1273 K 68

32 Microstructures in HERF Nitronic® 40 71

- 9 -10
 THIS PAGE

o
HANDBOOK OF DATA ON HYDROGEN COMPATIBILITY OF
STAINLESS STEELS

INTRODUCTION

This handbook is based on studies of hydrogen effects in

metals that have been conducted at the Savannah River Laboratory
(SRL) over the past 15 years. The major part of this continuing
study deals with austenitic, martensitic, and precipitation harden-
able stainless steels and is the subject of this handbook. The
handbook also includes data on high nickel alloys, iron-nickel,
iron-chromium alloys and pure nickel. These alloys are similar to
the stainless steels and their inclusion provides a more complete
coverage of the iron-chromium-nickel ternary alloy system.

Stainless steels have been utilized in structures and equip-

ment for hydrogen storage, rocket engines, and petrochemical
processing. Temperatures range from cryogenic to several hundred
Kelvin and effective hydrogen pressures often exceed 20 MPa. In
addition, hydrogen may be encountered as a corrosion product and
may contribute to degradation of equipment in such circumstances.
The data presented in this handbook covers the effects of hydrogen
on stainless steels at temperatures of 78 to 400 K and at pressures
up to 69 MPa. Therefore, the results are applicable to only a
portion of all possible industrial applications of stainless steel
in hydrogen environments.

There are two principal sections to this handbook: Alloy Data

Sheets which tabulate mechanical properties for each alloy and
Hydrogen Damage which discusses the hydrogen compatibility of
stainless steels. The mechanical property data for each alloy are
tabulated separately on one or more data sheets and indexed by
alloy number or trade name. The nominal compositions and mechani-
cal properties of the alloys are included in tables immediately
following the index. The discussion of hydrogen damage draws upon
published information from the technical literature to supplement
the SRL data.

The scope of the data accumulated on hydrogen damage varies

widely among the alloys that have been studied. In all cases,
tensile property measurements have been emphasized and fracture
modes examined. Fracture-mechanics-type specimens have been
studied in several alloys. Studies of Type 304L stainless steel
have been the most extensive. They include bar, cross-rolled
plate, and High Energy Rate Forged (HERF) stock. The effects of
grain size, heat treating temperatures, and quench rates on sus-
ceptibility to hydrogen damage have also been examined.

- 11 -
 Mechanical behavior, fracture modes, helium embrittlement
which may arise by beta decay of tritium, and hydrogen damage
mechanisms are discussed. Although a uniform and consistent data
base does not exist for each of the alloys, general trends in sus-
ceptibility to hydrogen damage have been identified. Changes in
mechanical behavior and fracture mode attributable to hydrogen have
been correlated with alloy composition, heat treatment, stress
state, and test temperature. The data base does provide sufficient
information for guidance in alloy selection and development of new
alloys tailored to specific service requirements.

During the past decade, several international conferences and

symposia have concentrated on hydrogen effects on metals. Refer-
ences 1 through 11. Proceedings of these meetings provide inter-
pretation and evaluation of hydrogen damage studies as well as
mechanical and physical property data on a variety of alloys other
than stainless steel. Hydrogen Damage, edited by C. D. Beachem,
Reference 12, is a selection of papers from the technical litera-
ture that provides an overview and insight into hydrogen damage, as
does the review paper by Hirth and Johnson, Reference 13.

Many of the publications and presentations at technical meet-

ings by researchers at the Savannah River Laboratory deal with
hydrogen effects in stainless steel. Only a few have been refer-
enced directly in the text, the remainder are listed in references
46 through 71. These papers supplement this Handbook and provide
more detailed discussion of specific topics.

- 12 -
SUMMARY

The susceptibility of stainless steels to hydrogen damage has

been evaluated primarily under rising uniaxial load rather than
under static or fatigue loading conditions. Smooth, notched, and
fatigue cracked specimens have been tested and hydrogen damage as-
sessed by measuring either fracture strain or the J-integral. As
shown in Table 1, stainless steels may be ranked qualitatively ac-
cording to the severity of damage measured on specimens precharged
with hydrogen at hydrogen pressures up to 69 MPa prior to testing.

Ranking of alloys will not in general be the same for static

or fatigue loading as for rising load conditions as seen in Table 2.
Slow-crack growth under static loading has been reported for A-286,
JBK-75, and Nitronic®-40 at stresses substantially below the ulti-
mate tensile strength, but not for Type 304L stainless steel.

The J-integral analysis of tensile data has been applied to

tests of smooth and C-shaped specimens. Both J at maximum load and
tearing resistance (change of J with crack length) are sensitive to
alloy composition, crack orientation relative to deformation tex-
ture, and test environment. The susceptibility to hydrogen damage
of several alloys was evaluated by the J-integral. Results were
similar to evaluation based on reduction- in-area in smooth bar
tensile test but more sensitive to orientation. The J-integral is
more versatile than a ductility parameter as it is applicable to
specimen designs that do not have an easily measured ductility
parameter.

Helium embrittlement at ambient temperature has been found in

five types of austenitic stainless steel: HERF Type 340L, HERF
Type 316, HERF Nitronic® 40 (21-6-9 steel), HERF A-286, and an-
nealed JBK-75 (a modified A-286). Helium was generated within the
specimens by radioactive decay of tritium which had been diffused
into the alloys at 420 K from high pressure (65 MPa) gas. Fracture
toughness of the alloys decreased with increasing accumulation of
helium. Damage was least in HERF Type 316 and greatest in HERF
Nitronic® 40 and JBK-75. Sustained load tests indicate that the
stress intensity for crack initiation and propagation also de-
creases with increasing helium concentration. Helium damage ap-
pears to be additive to the damage caused by the hydrogen isotopes
and on a per atom basis appears to be more severe than hydrogen
damage.

Additional observations on hydrogen damage may be derived from

the data:

- 13 -
TABLE 1

Ranking of Hydrogen Damage*

Minimal Moderate Severe .

316 304N CG-27
310 304L 17-4
Inconel® 718 216
309S** Nitronic®-40 Tenelon®
Nitronic®-50 A-286 18-18 Plus®
Incoloy® 800H JBK-75 Ni
XI8-3 Mn** AM363** AM350
N'i-SPAN-C** 18-2 Mn**
Carpenter 20Cb--3

* Based on tensile ductility. Nominal compositions

of these alloys are listed in Table AI.
** Ranking of these alloys is tentative.

TABLE 2

Hydrogen Damage Susceptibility

Property or Relative Property*

Type Type A-286 Nitronic®-
316 304L (JBK-75) 40
Tensile Properties
(Rising Load)
Notch Strength (H2/He) - 0.74 0.86 0.72
Notch Ductility (H2/He) - 0.50 0.38 0.19
Disc Burst Pressure (H2/He) 1.00 0.55 0.98 —
(Rising Load)
Fatigue Life (H2/He) 1.00 — — —
(Fluctuating Load)
Slow Crack Growth
(Sustained Load)
Threshold Stress No Crack No Crack 44-116** 100
Intensity Growth Growth MPa Vm~ MPa VS"
Detected Detected
in HERF 100 MPa Vm"
132 MPa

* Values <1.00 show effect of hydrogen by comparison with data

obtained in helium atmospheres.
** Variable depending upon heat treatment and precipitate morphology.

- 14 -
In Fe-Cr-Ni base alloys environmental hydrogen compatibility at
room temperature is maximized for alloys containing 15 to 30%
nickel. The role of other alloy elements has not been estab-
lished, but small additions of molybdenum (2 to 3%) appear
beneficial.

In Fe-Cr-Ni-Mn alloys, hydrogen compatibility improves as nickel

content increases. Manganese is not a substitute for nickel for
hydrogen compatibility in spite of the fact that manganese like
nickel stabilizes austenite.

There is a ductility minimum in the temperature range 200-300 K

where hydrogen damage is most pronounced. This ductility mini-
mum is observed even in alloys which show minimal hydrogen
damage and no fracture mode change such as Type 310 and 316
stainless steels.

Nitrogen contents greater than about 0.3 wt % appear to aggra-

vate hydrogen damage in Fe-Cr-Ni-Mn alloys.

Impurities normally present in commercial austenitic stainless

steels (silicon, manganese, sulfur, and phosphorus) do not
appear to influence susceptibility to hydrogen damage relative
to pure alloys. Segregation of these elements is potentially
harmful, however.

Increased strength with no loss or a small improvement in hydro-

gen compatibility is achieved by high energy rate forging (HERF)
of austenitic stainless steels.

The presence of hydrogen often evokes brittle fracture modes

under loading and temperature conditions where ductile failure
by microvoid coalescence ordinarily occurs. However, fractogra-
phy by itself is not adequate to diagnose hydrogen assisted
fracture in stainless steels because each brittle fracture mode
has been observed under other conditions.

- 15 -
HYDROGEN DAMAGE

Hydrogen damage occurs to all structural alloys, including

stainless steels.^^"^^ The widespread occurrence of hydrogen
damage is directly connected to the easy absorption of hydrogen
into metals and the high mobility of hydrogen in metals. Hydrogen
interacts with lattice defects, impurities, and internal boundaries
leading to a nonuniform distribution even for the equilibrium
state. Consequently, hydrogen is segregated along boundaries and
in regions of high stress or strain, thus localizing hydrogen
damage in regions which are often favorable sites for microcrack
initiation. Hydrogen segregation is more pronounced in ferritic
than in austenitic steels because of the larger elastic strain of
the hydrogen atom in the lattice of the ferritic steels.^^

There is no one physical process that is responsible for

hydrogen embrittlement of metals.^^~^^ At least three distin-
guishable forms of hydrogen degradation are recognized: hydride
embrittlement, hydrogen attack, and hydrogen-assisted fracture
(HAF). The first two processes are understood reasonably well,
whereas the last process has not yet been rationalized satisfacto-
rily. Hydride embrittlement occurs in metals such as niobium,
titanium, and zirconium, which form distinctive hydride phases
which are thermodynamically stable under ordinary ambient condi-
tions. These hydrides have a larger specific volume than the base
metal from which they form and appear to be inherently brittle.
The hydrides form structures with low crystal sjnmnetry. Brittle-
ness arises because of the limited number of slip systems available
in these crystal structures.

Hydrogen attack applies to formation of a gaseous product by

reaction of dissolved hydrogen with constituents such as oxygen and
carbon or with oxides and carbides present in the alloy. At ele-
vated temperature, steam or methane forms and generates microcracks
containing the reaction gas under pressure. This form of hydrogen
damage has been found in copper (steam embrittlement) and carbon
steels (methane formation). The addition of nitrogen to stainless
steels strengthens the austenite and creates a situation where
reaction between hydrogen and nitrogen could produce ammonia and
cause damage in the same way as oxygen or carbon. This form of
hydrogen attack has not been reported, however.

Hydrogen-assisted fracture occurs in many metals, including

the stainless steels. This form of hydrogen damage is character-
ized by the occurrence of a ductility minimum, an inverse strain

- 16 -
rate sensitivity, and crack propagation under static load at a
stress intensity less than ;the fracture toughness of the alloy.
The nature and degree of the damage vary widely, however, depending
on alloy composition, temperature, loading mode (static, cyclic,
increasing), stress state, and the environment. Only a few valid
generalizations can be made for hydrogen-assisted fracture of
stainless steels because the above influencing factors have not
been systematically investigated.

There is experimental evidence that internal boundaries may be

needed for significant hydrogen damage to occur. Investigations of
pure iron have demonstrated that the strength and fracture mode of
single crystals are unaffected by hydrogen either precharged to the
metal cathodically or charged during straining.^ Likewise, frac-
ture of single crystals of pure nickel is unaffected by hydrogen.^
It is to be noted, however, that serrated yielding was observed
during tensile straining of nickel at temperatures of 150 to 270 K,
providing evidence of hydrogen binding to dislocations. In con-
trast, polycrystalline specimens of iron and nickel are known to be
severely embrittled by hydrogen. No hydrogen damage studies have
been attempted with single crystals of any stainless steel. How-
ever, crystals free of internal boundaries may not be attainable
because of the ease of formation of annealing twins in these steels.
As noted later in this report, hydrogen-assisted fracture of stain-
less steel frequently propagates along twin or grain boundaries.
Furthermore, ductile fracture processes can be altered by hydrogen
accumulation at inclusion-matrix interfaces by diffusion or a
dislocation transport mechanism.1" Hydrogen may alter either
nucleation or growth of microvoids, depending upon the nature of
the inclusion-matrix or precipitate-matrix interfaces.'^^ Conse-
quently, dimple size on fracture surfaces may be either decreased or
increased by the presence of hydrogen.

MECHANICAL BEHAVIOR

Hydrogen has a pronounced effect on crack initiation and growth

but only a small effect on deformation behavior. Therefore, mechan-
ical properties that are sensitive to crack behavior are suited to
hydrogen damage studies. The most severe tests involve high stress
or large plastic strain. Fatigue and creep rupture are two such
tests. However, tensile tests of smooth and notched round speci-
mens have been the principal methods for evaluating the suscepti-
bility of stainless steels to hydrogen damage because of low cost
and simple test procedure. Change in reduction-in-area has been
the most commonly used index of hydrogen damage. Generally, if
there is a loss in reduction-in-area, other mechanical properties
are affected also; however, the converse is not true. Consequently,
the tensile test is of limited applicability as a screening test
for susceptibility to hydrogen damage. Recently, J-integral

- 17 -
techniques have been applied to analysis of tensile data of fatigue
precracked specimens, either single-edge notched (SENT) or C-shaped.
Impact specimens and internally pressurized tensile tubes have been
tested in limited numbers.

Specimens have been tested following exposure to high-pressure

(about 1 MPa to 69 MPa) hydrogen (internal hydrogen damage) or in
the presence of high-pressure (69 MPa) hydrogen gas (external hy-
drogen damage). Tests with an external hydrogen source were con-
ducted at room temperature only. In contrast, mechanical testing
of hydrogen-charged specimens has been done at 4.2 and 78 K and at
temperatures between 200 and 400 K.

Mechanical property data for the stainless steels that have

been studied at SRL are collected in the Alloy Data Sheets. The
tables are organized by alloy and by type of test for each alloy.

Material and Processing Variables Affecting Hydrogen Damage

There are distinct differences in the degree of hydrogen

damage among the stainless steels. Alloy composition provides a
simple scheme for classifying steels according to severity of hy-
drogen damage. Such a criterion does not, however, provide any
insight into the reasons for the observed differences in response
to hydrogen. Properties such as phase stability, stacking fault
energy, or dislocation slip mode are related to composition and are
directly related to the processes of crack initiation and propaga-
tion. Microstructure as established by heat treatment and mechani-
cal processing modifies hydrogen damage for any given composition
of steel. Thermal and mechanical treatments influence severity of
hydrogen damage in a steel through changes in grain size, alloy and
impurity segregation, precipitate size distribution and coherence,
and dislocation substructures. Severity of hydrogen damage may be
altered in practice by alloy composition, heat treatment, and
mechanical processing. Examples of these effects are cited below.

Ccmposition

Stainless steels considered in this report fall into two broad

composition classes: iron-chromium-nickel and iron-chromivim-
nickel-manganese. The latter class of alloys all contain nitrogen
which is added to strengthen the austenite.

Environmental embrittlement of iron-chromium-nickel alloys

correlates with nickel content, as seen in Figure 1. These results
were obtained from tensile tests in a 69 MPa hydrogen environment
at room temperature. This correlation appears valid for both high
purity and commercial grades of steel. Consequently, elements

- 18 -
 100

80
% Conmercial Alloys
A High-Parity Alloys

i
o
a3 60

c
0)
E-
40 —
C

20 —

0 1! 1
10 20 30 40 50 60
Nickel. wt,%

FIGURE 1. Enviroiunental Hydrogen Damage in Fe-Cr-Ni

Alloys at Room Temperature and 69 MPa Hydrogen

- 19 -
present as minor alloy constituents such as manganese, silicon,
phosphorus, and sulfur are not primary sources of hydrogen damage.
As will be shown later, however, segregation of some minor alloying
elements can alter the extent of hydrogen damage.

The composition range from 8 to 14% nickel, where the marked

improvement in resistance to hydrogen damage occurs, corresponds to
increased austenite stability and decreased formation of martensite
during plastic deformation. According to the equilibrium diagram
calculated by Breedis and Kaufman for pure iron-chromium-nickel
alloys, austenite is the stable phase above 19% nickel for a 20%
chromium alloy at 300 K.'° The high-purity alloys tested at SRL
indicate that as little as 14% nickel will stabilize austenite with
respect to transformation to cx'-martensite at room temperature.
Furthermore, austenite stability with respect to transformation to
the epsilon phase increases as the nickel content is increased from
10 to 20%. Figure 2 shows the composition dependence of stacking
fault energy, which is a measure of austenite stability with re-
spect to the epsilon phase,'^^ Both a'-martensite and epsilon phase
are detrimental, as resistance to hydrogen damage is rapidly im-
proved over the range of nickel concentrations (10 to 20%) corre-
sponding to increased austenite stability. However, as discussed
later, the relation between martensite and hydrogen damage is
neither simple nor clearly demonstrated.

Hydrogen damage in the Fe-Cr-Ni-Mn alloys as measured by

ductility in high-pressure hydrogen does not correlate uniquely
with nickel or with equivalent nickel [Ni + 30 (C + N) + 0 . 5 M n ] .
For several compositions, Figure 3, a range of results may be
obtained. This range of behavior is probably caused by variations
in nitrogen content as documented for Nitronic® 40 (21-6-9 stain-
less steel). The higher nitrogen contents not only strengthen
the alloy but also reduce the stacking fault energy, thereby promot
ing coplanar dislocation arrays and formation of epsilon phase.^^

The influence of other major elements such as chromium, man-

ganese, and molybdenum on susceptibility to hydrogen damage is not
well documented. Most commercial stainless steels contain 15 to
25% chromium with 18-20% being the most common concentration. Var-
iations in hydrogen damage among the stainless steels cannot be
linked directly to chromium because the other alloying elements
change simultaneously. Chromium stabilizes ferrite, which is more
susceptible to hydrogen damage than austenite, therefore, a high
chromium content is expected to be detrimental.

Molybdenum appears to alleviate hydrogen damage when present

in small amounts, 2-3%, as indicated by comparison of Type 316 with
Type 304 stainless steel. However, Type 316 stainless steel also
has a nickel content about 2% higher than Type 304 stainless steel.
The higher nickel cannot account completely for the reduced

- 20 -
 Cr

, _ - ^ ^ ^ , 5 5 . ., -

- -CO

V >E>«rg}fi*i •tiij/m^/ :-;^-\

•;-;v~^-^-3';-^

ii:::-:^
-:^

.-.•.•/.-.'.•.•/.•.•.•.•/.VN
:\

-.rf'-
^^•A^^^V.^^^^^^^^•^^^^^^^^^^^^^^^^^^^•/•.^^>^^^^•/^^^^yrA^y^Vv

•^.vV-v^x
^ ?fe':X:::yf
\v'-"-"-^>"-'-'-Vr'-'-^*%'-"--^y-'-"Ay~v.^y-'--':v/-^"<;7:^.-%\7^^^
Fe

FIGURE 2. Stacking Fault Energies in Fe-Cr-Ni Alloys

at Room Temperature (Ref. 19)

- 21 -
 1 1
100 9 <> —
© 9

—
B-« '^

>r 80 •-

O 18-18 Plus®
a "" • Tenelon®
3
©18-3 —
60 —
©216
•H
CO
—
) €> Nitronic®-40 —
C «
OJ « Nitronic®-50
H

c 0
•H
CO —
(1)

20 —

1 1 1
10 15 20
Nickel, wt %

FIGURE 3. Environmental Hydrogea Damage in Fe-Cr-Ni-Mn

Alloys at Room Temperature and 69 MPa Hydrogen

- 22
hydrogen damage because an alloy with 14% nickel with no molybdenum
shows greater hydrogen damage than Type 316 stainless steel. Large
concentrations of molybdenum (>5%) are expected to aggravate hydro-
gen damage because molybdenum stabilizes ferrite.

Manganese stabilizes austenite and should improve hydrogen

performance of stainless steels. However, alloys with high man-
ganese contents and little or no nickel such as Tenelon® (U. S.
Steel Corp.) and 18-18 Plus® (Carpenter Technology) behave very
poorly in hydrogen, Figure 3, indicating that austenite stability
in itself is not sufficient to minimize hydrogen damage.

Heat Treatment

Heat treating has been shown to affect susceptibility to

hydrogen damage in several ways: change of grain size, segregation
or redistribution of impurities and alloying elements, and precipi-
tation of new phases. The general patterns of variation in hydro-
gen damage with composition can be modified, therefore, by heat
treatment.

The only study of the effect of grain size on hydrogen damage

has been made with Type 304L stainless steel. As seen in Figure 4,
flow stress was increased by finer grain size and the effect was
greater in specimens that were saturated with hydrogen. Grain size
effect on the yield strength of specimens tested in high-pressure
hydrogen, was the same as for tests in helium, Figure 5. Presuma-
bly, comparable results would be obtained with the other alloys.

The grain size dependence of the flow stress (^f) follows

the Hall-Petch relation, ^f = ^Q •*• kfd~^'2, where o^ and kf are
functions of strain and d is grain size. Analysis of the data
demonstrates that hydrogen strengthening is associated with the
back stress of dislocation pileups against internal boundaries and
the increased resistance to dislocation intersection where hydrogen
is present.2^ Hydrogen causes only a small increase in lattice
friction stress in keeping with the relatively small lattice dis-
tortion caused by an interstitial hydrogen atom in the face cen-
tered cubic lattice.^•'^

Sensitization of austenitic stainless steels occurs during

heating to temperatures in the range 720 K to 1150 K, or upon cool-
ing slowly through this range. Chromium-rich carbides of the gen-
eral form M23C5, where M is chromium or iron, precipitate along the
grain boundaries and deplete the boundary region of chromium. In
nitrogen strengthened steels, the carbides may contain nitrogen
also.

- 23 -
 1000

O No Hydrogen
• Hydrogen Charged

Q.

500
0)

0 ±
0.1 0.2
I
0.3 0.4
(Grain Size)""^. {fjim)~^^^

FIGURE 4 . Hall-Petch P l o t for Tensile-Text Specimens

of Type 304L S t a i n l e s s S t e e l at a P l a s t i c
Strain E - 0.05

- 24 -
 400

100 —
O Tested in Helium Environment
• Tested in Hydrogen Environment

0.1 0.2 0.3 0.4

1/2
Grain Diameter, (/im)

FIGURE 5. Hall-Petch Plot of Yield Strength for Type 304L

Stainless Steel Tested in in High-Pressure Helium
or Hydrogen Environments

- 25 -
 The effects of sensitization anneals at about 920 K on hydro-
gen damage have been reported for Type 304,^^ 304L, 309S,^^ and
Nitronic® 40 stainless steels. These investigations of hydrogen-
assisted fracture of sensitized austenitic stainless steel do not
present a consistent or uniform set of mechanical property data.
Consequently, comparison of relative strength or ductility changes
due to sensitization are limited. In those cases where data were
collected, sensitization appears to lower the tensile strength of
the alloys both for tests in air and in hydrogen. The ductility of
sensitized specimens in hydrogen was less than that of the solution
annealed specimens except for Type 309S steel, where no ductility
change or intergranular fracture was observed. The anomalous be-
havior of sensitized Type 309S steel is not attributable to compo-
sition effects such as carbon/chromium ratio or carbon + nitrogen/
chromium ratio. Carbide precipitation was observed in all four
alloys.

Hydrogen compatibility of precipitation hardenable alloys is

sensitive to aging temperature and time. The fracture toughness of
single-edge notched specimens of 17-4PH stainless steel varied sub-
stantially with aging. The most severe degradation occurred for
the peak aged specimens. In the case of A-286, aging times of 4,
8, and 16 hours at 993 K resulted in small changes in fracture
toughness and improved hydrogen compatibility for the shorter aging
t imes.

Mechanical Processing

Parts for hydrogen service may be made from annealed or cold-

worked bar and cross-rolled plate, or HERF stock. Further cold-
working may occur during fabrication. The various working proc-
esses introduce characteristic microstructural features, affect the
mechanical properties, and may change susceptibility to hydrogen
damage of the stainless steel.

Suceptibility to hydrogen damage observed in tensile tubes

machined from cross-rolled plate was unchanged by prior cold-work
of up to 25%, Figures 6 and 7. Specimens were filled with hydrogen
at 69 MPa pressure and stored 30 days at 430 K before tensile test-
ing at room temperature. The bore of the specimen contained high-
pressure hydrogen during testing.

In the case of thin sheet (0.1 mm thick), a progressive

decrease in tensile ductility was found as prestrain was increased
from 0 to about 12%, Figure 7.

In contrast, properties of annealed and 10% cold-worked sheet

of Type 309S stainless steel were comparable. After three and a
half months' exposure to hydrogen at 47.5 MPa at 345 K, tensile

- 26 -
 1000

10
a.

286 mol Hz/m^

200 h O Test in Air

• 69 MPaHa, 430K, 30 days
0 I L
0 10 20 30
PRESTRAIN, %

FIGURE 6. Effect of Prestrain on Tensile Strength of Type 304L

Stainless Steel Tubes in Air' or Hydrogen

- 27 -
 T T
• n Sheet - 0.25 mm thick, 69 MPa,
345 K, 3-1/2 months

A Sheet - 0.1 mm thick, 69 MPa,

345 K, 43 days

• o Tensile Tube
69 MPa, 430 K, 30 days
Open - In Air

Filled - H2 Charged. Numbers on

curves are estimated H2
concentrations at specimen
surface.

188 mol Hz/m'

10 20 30
PRESTRAIN, %

FIGURE 7. Effect of Prestrain on Elongation of Type 304L

Stainless Steel in Air or Hydrogen

- 28 -
strength was 625 MPa for both alloys, and total elongations were
51 to 55% for the annealed sheet and 48 to 57% for the cold-worked
sheet. Properties of annealed sheet were a 570 MPa tensile
strength and 59 to 64% elongation prior to hydrogen exposure.

A British study on two stainless steel compositions showed

that cold work increased susceptibility to hydrogen damage,^5
Figure 8. These steels have no counterpart among commercial stain-
less steels in the United Stated. The results further demonstrate
that the effect of cold-working is greater in the 12Cr-12Ni alloy
than in the 23Cr-2lNi alloy. Martensite is formed during cold-
working of the former alloy, but not the latter alloy, and may
account for their difference in behavior.

Tensile test data show that HERF alloys are more resistant to
environmental hydrogen damage than either annealed or cold-worked
alloys. Table 3. A similar result is observed for thermally
charged specimens tested at room temperature. At lower test tem-
peratures, the HERF specimens of Type 304L and Nitronic® 40 stain-
less steels are only marginally better than annealed specimens in
terms of ductility. Yield and tensile strengths are higher in the
HERF alloys, therefore, on the basis of retained ductility for a
given strength level, the HERF alloys are superior to annealed
alloys.

Specimens made of Nitronic® 40 cold worked 30% from HERF stock

were slightly less ductile than specimens in the HERF condition
during tests in high-pressure hydrogen. Figure 9. Tensile data
also suggest a slight increase in susceptibility to hydrogen
damage , Figure 10.

HERF develops a distinctive deformation pattern in stainless

steels which has its origins in segregation in the ingot, and
rolling of the plate before HERF. The effect of this pattern on
susceptibility to hydrogen damage was evaluated by the J-integral.
J^, the critical force at maximum load (area under load deflection
curve), and tearing resistance dj/da (change in J with crack
length, a, at maximum load), were evaluated for V-notch specimens
of Type 304L, A-286, Nitronic® 50, and Nitronic® 40 stainless
steels machined from HERF bars. Notches were oriented as shown in
Figure 11 with respect to the forging pattern and fatigue pre-
cracked at a maximum stress intensity of 44 MPaVmT Specimens were
tested in high-pressure helium or hydrogen and after prior exposure
to deuterium at 69 MPa pressure for seven days at 520 K or three
weeks at 620 K.

The deformation pattern caused by HERF processing affected

J^ but not dj/da for tests in high-pressure hydrogen or helium.
The lowest values of J^ were for notches parallel to the flow
lines, an effect especially pronounced in A-286 and Nitronic® 50

- 29 -
 Hydrogen Damage Susceptibility
Change in Percentage Reduction in
8 Area per 1 cm' Hydrogen per 100 gram alloy

00

M H
C hh
« ht»
n n
n n
•o rr
ft

•-• n
v< O
1
H» O
O O
CO i->
ft O ,

n K
«
» o
ft pd
» *<
n o.
I- 1
-.gn
o

OQ
TABLE 3

Mechanical Properties of HERF and Annealed Alloys in High-Pressure Gas

Test* Strength, MPa** Ductil ity %

Alloy Condition Environment Yield Ultimate Elong Ra
304L Annealed He 186 565 74 81
H2 206 503 48 33
HERF He 490 660 43 77
H2 490 670 41 69
310 Annealed He 179 483 67 82
H2 186 486 66 79
HERF He 470 606 21 71
H2 460 620 24 70
Nitronic® 40 Annealed He 352 689 58 78
H2 358 696 59 77
HERF He 570 780 34 75
H2 570 790 30 73
A-286 Annealed He 724 1117 26 47
H2 710 1131 34 49
HERF Air 875 1060 24 29
H2 834 1110 27 32
CG27 Annealed He 806 1165 29 26
H2 855 1117 10 12
HERF He 1070 1385 12 12
H2 1034 1138 1 3

* He = 69 MPa He; H2 = 69 MPa H2; Air = 0.1 MPa air,

** Yield = True stress at 5% strain.
Ultimate = True stress at maximum load.

- 31 -
 60 —

AIR
&«
40 — Hz - 69MPa

CD

20

ANNEALED HERF HERF + 30% CW

EXPOSED
188 mol Hz/m^

FIGURE 9. Effect of Processing on Hydrogen Damage in

Nitronic* 40 Elongatibn

- 32 -
 AIR— 1 — H2 — 69MPa
200
—1
000 — -

800
ZIZ
h-
CD
2:
LU 600 — -
OH
I—
00

400 — —

200 — —

n
ANNEALED HERF HERF + 30% CW EXPOSED
188 mol Hz/m^

FIGURE 10. Effect of Processing on Hydrogen Damage in

Nitronic® 40 Strength

- 33 -
A. Cross Section of Bar B. Parallel Orientation of Notch
Showing Forging Flow Lines

C. 45° Orientation of Notch D. 90° Orientation of Notch

FIGURE 11. Orientation of V-Notch with Respect to Forging

Flow Lines in Tensile-Test Specimens Machined
from Bass of HERF Stainless Steels

- 34 -
stainless steel. However, hydrogen damage as measured by J^, was
greatest for notches 90° to the flow pattern except for Nitronic® 40,
where damage was greater for 0 and 45° orientations than for the
90° orientation. dJ/da was affected relatively less by hydrogen
damage than JJQ.

Changes in J^j and dJ/da caused by hydrogen charging prior

to testing were alloy and orientation dependent. Hydrogen exposure
of A-286 stainless steel decreased J,ji and dj/da. For Type 304L
stainless steel, dj/da was unaltered and J^i was decreased for 45°
and 90° orientations. For Nitronic® 40, J.^ and dj/da were in-
creased by hydrogen exposure prior to testing, possibly because of
microstructural changes that occurred during hydrogen charging.

Effect of Martensite

The above results show that the effect of cold work on hydro-
gen damage is variable and may be linked to martensite. A study
was made of the relation between martensite and hydrogen damage.
Martensite was formed in annealed specimens of Type 304L austenitic
stainless steel either by thermal or mechanical treatment. Speci-
men surfaces were electropolished to remove the deformed outer
layer that is generated during machining. This layer contains some
a'-martensite in Type 304L stainless steel. The volume fracture
martensite formed by the various treatments was not measured di-
rectly, but the relative quantity of martensite was indicated by
the magnetic response of the specimens. The specimens were pulled
to failure in a high-pressure (69 MPa) hydrogen atmosphere follow-
ing thermal or mechanical treatment. Thermal treatment consisted
of quenching specimens to 77 K in liquid nitrogen and holding at
temperature for 10 minutes. Three mechanical treatments were used:
cold reduction of 12 or 26% by cold swaging, torsional strain
(70-in. lb) at room temperature or 77 K, and 10% tensile strain at
room temperature or 77 K. Tensile strain at room temperature
caused some transformation to e-phase and a'-martensite formation.

Uniform elongation and plastic strain to failure for each of

the treatments are reported in Table 4. Magna-gage® (American In-
strument Company) readings following treatment and prior to tensile
testing are given. Several important observations are apparent
from inspection of the table: ductility is reduced when martensite
is present prior to testing; hydrogen embrittlement does not cor-
relate simply with volume fraction martensite, as the highest
Magna-gage® readings ( B ) do not correspond to lowest ductility ( F ) ;
the presence of an applied stress during the quench to 77 K sup-
presses transformation to a'-martensite (B, C, and D ) ; the radial
distribution of the martensite and possibly its orientation rela-
tive to the applied stress affect hydrogen embrittlement (B, C, F ) .

- 35 -
TABLE 4

Effect of Martensite on Subsequent Environmental Hydrogen Damage in

Type 304L Stainless Steel

Ductility
Uniform Fracture Magna-gage®
Treatment Elongation Strain Reading

A. Annealed-electropolished 79 1.31 0

B. Immersed in liquid nitrogen 51 0.45 90 100

for 10 min

Applied torque and immersed 58 0.38 30

in liquid nitrogen for 10 min

Applied tensile stress and 60 0.42 13

immersed in liquid nitrogen
for 10 min

10% plastic strain at room 75 1.11

temperture

7.9 Nm torque at room 26 0.24 30

temperature

- 36 -
Environmental Variables Affecting Hydrogen Damage

Degradation of mechanical properties by hydrogen depends upon

several factors external to the specimen: temperature, hydrogen
pressure, stress state, and strain rate. These factors control how
much hydrogen enters the material and how rapidly solution and
distribution of hydrogen occur within the specimen. Hydrogen re-
distribution to regions of high stress and the hydrogen diffusion
rate relative to any imposed deformation rate directly influence
severity of hydrogen damage.

Hydrogen Pressure

The severity of hydrogen damage in any alloy has been shown to

relate directly to hydrogen pressure whether for hydrogen-charged
specimens or for testing in a hydrogen environment. As seen in
Figure 12, tensile ductility of Type 304L stainless steel falls off
rapidly as hydrogen pressure is increased up to 14 MPa and then
more slowly up to 69 MPa. Tensile strength of notched and fatigue
cracked specimens also decreases from 780 to 660 MPa as the hydro-
gen pressure increases from 0.1 to 6.9 MPa. For most other alloys,
testing has been at fixed hydrogen pressure of 69 MPa, and data do
not exist for other pressures.

A pressure effect is seen in both tensile and threshold stress

intensity data for Inconel® 718.^° Tensile strength and
reduction-in-area decrease linearly with square root of hydrogen
pressure at pressure above about 10 MPa. The stress intensity at
crack arrest for wedge opening loaded specimens at room temperature
decreased from 90 MPa Vm"to about 42 MPa Vm for hydrogen pressures
of 20 MPa and higher.

Hydrogen gas pressures of 69 MPa have no significant effect on

the mechanical properties of Type 310 stainless steel. However,
severe hydrogen damage can be induced in Type 310 stainless steel
by cathodic charging where the effective hydrogen pressure is ex-
tremely high (»200 MPa).27

Test Temperature

Hydrogen damage in stainless steels has been shown to be tem-

perature dependent for alloys where appropriate testing has been
done. The most severe damage occurs at 220-280 K. As seen in
Figure 13, tensile ductility in several alloys is a minimum in this
temperature range. There does not appear to be any significant
difference between Fe-Cr-Ni and Fe-Cr-Ni-Mn alloys in this respect.
Figure 14. Furthermore, the minimum ductility ratio is observed in
this same temperature range even in alloys which are affected very

- 37 -
 4 6 8
Test Pressure, (MPa)^/^

FIGURE 1 2 . Effect of Hydrogen Pressure on Ductility of

Type 304L Stainless Steel Tested to Fracture
in Hydrogen at Room Temperature

- 38 -
 400
Test Temperature, K

FIGURE 13. Ductility Minima in Fe-Cr-Ni Alloys Charged with

Deuterium at 69 MPa and 620 R for Three Weeks

- 39 -
 u
•H 1.0

CM

X
0.8
OJ

3
rH
•H
0.6

O
•H
4J
cS
0.4

c
•H
cd
O Tenelon"^
0.2
o • Nitronic®-40
•H
w © Nitronic®-50
to
rH
P-l
I
100 200 300 400
Test Temperature, K

FIGURE 14. Ductility Minima in Fe-Cr-Mn-Ni Alloys Charged

with Deuterium at 69 MPa and 620 R for Three Weeks

- 40 -
little by hydrogen, such as Nitronic® 50, Type 310, and Type 316
stainless steels. It is not known yet if the threshold stress
intensity for sustained load crack growth in these stainless steels
follows a similar temperature dependence. However, the data for
Inconel® 718 (Huntington Alloys) lead to the inference that Kj-^j and
tensile ductility have a similar temperature dependence.2"

The general trends of both test temperature and nickel content

on hydrogen damage are illustrated in Figure 15 for three high-
purity alloys of constant chromium content (about 18%). Isoductil-
ity lines sketched in this figure outline a region of significant
hydrogen damage (£_ '^ 0.5) at about 200-240 K that extends to 14%
nickel.

Although most materials show a hydrogen-induced ductility min-

imum, the temperature of the minimum is not always around 220-280 K
as in stainless steels. Pure nickel displays ductility minima
around room temperature, at 220 K and at 150 K, but the 220 K
minimum is the only one to persist over a wide range of strain
rates (10"^ to 10"^ sec~^).^^ Maximum hydrogen damage in Inconel®
718 is at room temperature as measured by ductility, notched ten-
sile strength, and threshold stress intensity.2"

Strain-induced martensite forms during low-temperature defor-

mation of many austenitic stainless steels which are nominally
stable at room temperature. For a given alloy, there is a tempera-
ture designated M j , above which no martensite forms on deformation
and below which deformation produces martensite. The amount of
transformation per unit strain increases as the temperature is
lowered below M^j. Table 5 lists approximate M^j temperatures for
several of the austenitic steels that have been investigated. The
appearance of brittle fracture and degradation of mechanical prop-
erties does not correlate with formation of strain-induced martens-
ite. In particular. Type 316 stainless steel does not show a sig-
nificant change in fracture strength or fracture mode between 250 K
(no martensite) and 200 K (martensite present).

Magnetic measurements on specimens of Type 304L stainless steel

strained incrementally at 200 K show that hydrogen suppresses forma-
tion of strain-induced martensite. Figure 16. However, hydrogen
appears to enhance martensite formation during deformation at 250
and 298 K. Suppression of martensite formation was seen also in
Tenelon® at 200 K and Nitronic® 40 at 78 K, Figure 17.

Stress State and Strain Rate

Plastic deformation and fracture characteristics are affected

by stress state and strain rate. Severity of hydrogen damage.

- 41 -
 O No martensite
400
detected
• Martensite
present after
test.

300
Contours show
equal values of
plastic strain
Ep (ductility)
200

100

10 20
Nickel, wt %

FIGURE 15. Isoductility Diagram for

Hydrogen-Charged Fe-Cr-Ni Alloys

- 42 -
TABLE 5

Temperatures for Strain-Induced Martensite Formation

Alloy* Md Temp, K

304L 350
310 <4
316 220
I800H <78
Tenelon® 200*
Nitronic® 40 ~100
Nitronic® 50 ~150
A >370
B 250
C 150

* Martensite forms at temperatures near 200 K,

but not at either higher or lower temperatures.

- 43 -
 1000

900

800 ^Magna-Gage® Saturated

KOA)
O.700
c
S 600 Open — No D2
Closed — Saturated D,^
© D
S,500

S 400
(0

300- 198 K

200-

100-

0L_
0.0 0.3
• V • .^•V
0.4
298 K

V
0.5
t * I
f
0.6
Plastic Strain

FIGURE 16. Effect of Hydrogen on Strain-Induced Martenite

Formation in Type 304L Stainless Steel
400
O Nitronic®-40 78K
n Tenelon® 200K

300
Open No Hydrogen
Filled Saturated With
Hydrogen at 69
MPa, 620 K.
200

100 —

0.05 0.10 0.15 0.20 0.25

Plastic Strain

FIGURE 17. Effect of Hydrogen on Strain-Induced Martenite

Formation in Tenelon® and Nitronic* 40

- 45 -
therefore, is also affected by stress state and strain rate. Dif-
ferent test methods may assess hydrogen degradation of mechanical
properties differently, depending upon how stress state and strain
rate influence hydrogen interactions with the alloys.

A biaxial stress state occurs during burst testing of discs.

Hydrogen environment damage may be assessed by pressurizing discs
of the same material with helium or hydrogen gas and measuring the
burst pressure.28 A ratio of burst pressures of one indicates no
hydrogen damage, whereas ratios greater than one indicate hydrogen
damage with severity of damage increasing as the ratio increases.
Steels such as A-286, Type 310, and Type 316 are relatively immune
to hydrogen damage in this test; CG-27 behaves poorly; and Type 304
or 304L are variable, depending upon the treatment prior to testing.
The ranking of the alloys in the disc burst test is generally the
same as for notch tensile tests in helium and hydrogen for a pres-
sure of 69 MPa.

Notch Tensile Tests. Stress concentration and plastic

constraint at the root of a notch can increase the severity of
hydrogen damage under some circumstances. Notched specimens of
Type 304L and Nitronic® 40 stainless steels tested in high-pressure
hydrogen at room temperature illustrate the notch effect. The
strength and ductility of the notched specimens are degraded more
than that of the smooth specimens. Figures 18 and 19. In contrast,
ductility loss in notched specimens of hydrogen-charged Type 304L
stainless steel was less than in smooth specimens. Table 6. How-
ever, absolute ductility of the notched specimens is less than for
smooth specimens. Hydrogen charged specimens of Nitronic® 40 were
more severely damaged when notched than when smooth, as was the
case for tests in a hydrogen environment.

J-Integral Analysis. C-shaped and notched specimens were

fatigue cracked and then tested in hydrogen or helium at 69 MPa
pressure. Susceptibility to environmental and internal hydrogen
damage was evaluated by comparing J-integral values at maximum load
(J^j) and dJ/da for four stainless steels machined from HERF bar
stock. Ranking of the alloys was not consistent and depend on
which specimen orientation was chosen for the comparison (Table 7).

Based on plastic strain to failure (reduction in area), HERF

Type 304L stainless steel is more susceptible to environmental
hydrogen damage than HERF Nitronic® 40 stainless steel in direct
opposition to the ranking based on Jm. These measures of hydro-
gen damage are biased differently, however, with respect to defor-
mation and fracture. JQ| depends on flow stress and plastic strain
up to the point where crack advance is assumed to begin. Plastic
strain to failure, on the other hand, is affected strongly by the
ability of the alloy to neck down without breaking and is effec-
tively sampling a region of deformation not included in J^,.

- 46 -
Sensitized, Sensitized, Sensitized, Solution Annealed,
No Notch Shallow Notch Deep Notch No Notch

FIGURE 18. Relative Effect of Hydrogen on Properties

of Sensitized Type 304L Stainless Steel
 i./.

^f

^u ^u
1.0
Q)
X
M-l •-S
Q)
(.) EC
•—s
CM
>w<'
3
^f
X 0.8
&
(i)
M-l /--N
CM
^u
ts
S--'
0) «\ 3
n h
3
iH .t 0.6
•H o
to •H
fL, 4-1

4-1 ^
to
.C
o 4-1
•H
4-1
60 0.4
c
0)
^ 4J
CO
c
•H
to (1)
U 4-1
^f
4-1
V3
tfl n.2
a
•rl
1

O 4-1
•H iH
U P
tn
to •a
iH
P- c
to 0
Solution Annealed, Sensitized, Sensitized,
No Notch No Notch Notch

FIGURE 19. Environmental Hydrogen Damage in Nitronic* 40

Stainless Steel at Room Temperature

- 48 -
TABLE 6

Effect of Hydrogen Charging on Tensile Properties of Type 304L Stainless Steel

Plastic
Nominal T e n s i l e Strain
Condition Specimen S t r e n g t h , MPa To F a i l u r e

As received Smooth 600 1.50

Notch 770 0.30
Annealed* Smooth 600 1.43
Notch 710 0.24
Hydrogen charged** Smooth 530 0.37
Notch 580 0.13

* Annealed 200 days at 380 K in argon.

** Exposed to hydrogen gas at 69 MPa for 200 days at 380 K.

TABLE 7

Ranking of Hydrogen Compatibility (HERF Stainless Steels)

Environmental Hydrogen
Tensile*
dJ/da Ductility
'm

A-286** A-286 Nitr:onic® 40

N i t r o n i c ® 40**
304 L Nitronic® 50** 304 L
Nitronic® 40
Nitronic® 50** 304L

Internal Hydrogen
Tensile*
T **
'Jm dJ/da Ductility

Nitronic® 40
Nitronic® 50 N i t r o n i c ® 50 Nitronic® 40
N i t r o n i c ® 40
304 L 304 L 304 L
A-286 A-286

* Orientation dependence not measured.

** Ranking is orientation dependent.

- 49 -
Furthermore, compact tensile specimens such as the C-shaped speci-
mens of this study are not susceptible to unstable plastic tearing,
whereas tensile specimens can be driven to tearing instability by
internal microcracks in the necked region.

There is agreement in ranking of HERF Type 304L and HERF

Nitronic® 40 stainless steels between tensile ductility and tearing
resistance, dJ/da. In both cases, HERF Nitronic® 40 stainless
steel is less susceptible to hydrogen damage. In this comparison,
both parameters are influenced by the deformation and crack growth
processes that succeed the onset of crack growth at load maximum.

Static Loading. Static loading may lead to slow-crack

growth in high-pressure hydrogen in some stainless steels.
Figure 20.2°'^^"^^ The only alloys where crack growth was not
seen had yield strengths below 800 MPa. The specimens were too
thin to permit reaching a stress intensity high enough to cause
crack growth. In all cases, the threshold stress intensities are
substantially less than the fracture toughness. The scatter band
for fracture toughness derives from valid J-integral tests made at
4.2 or 77 K. KJ^Q may be higher or lower at room temperature
depending upon the steel.-*^~3^

Crack propagation has been observed in notched specimens of

Type 304 stainless steel at stresses equal to 80% of the ultimate
tensile strength.^^ These tests were in dry hydrogen at one atmos-
phere pressure at room temperature. A notched tensile specimen of
Type 304L stainless steel failed after 41 hours in air under a
stress of 490 MPa, or 85 percent of the notch tensile strength.
This specimen had been exposed to D/T for 200 days at 69 MPa and
380 K. An unexposed control specimen survived 200 hours at 650 MPa
(85% of notch tensile strength) without failure. In another ex-
periment with smooth specimens of Type 304L stainless steel, no
crack growth was observed after cathodic charging for stresses
close to the ultimate tensile strength.-'" As seen in Table 8, a
pre-existing crack did not progagate in Type 304L stainless steel
tensile tube until the stress reached a very high level. Also,
crack growth has not been observed in Type 310 stainless steel at
80% of the ultimate strength.^^ In A-286 stainless steel, hydrogen
enhanced crack growth was observed at stresses above the yield
strength, but tensile properties were unaffected by exposure to
69 MPa hydrogen.31 The threshold stress intensity for crack growth
in 35 MPa hydrogen was less than 113 MPa Vm" compared to greater
than 154 MPa Vm" in helium.

Measurements of sustained-load cracking were made on a modified

A-286 stainless steel (JBK-75) in hydrogen gas at 100 and 200 MPa
pressure.32 yield strengths varied from 690 to 930 MPa with dif-
ferent thermomechanical treatments. Crack advance was intergranular

- 50 -
500
T A K-TH in 100-200 MPa H,
\ (Reference 32)
O K-j-H in 35 MPa H j
(Reference 26)
400
T No Crack Propagation
\ \

\ \
Fracture Toughness Data for
300 Stainless Steel at 4 to 77K
(Reference 34-35)

\ \

200 - \ \

\
Nitronic \
316^50 ^
\

304
LI AJBK-7&
100 - A286A^A316
Nitronic-40

i
O £^ J B K - 7 5 \
625
903
A
304 L JBK-75 z9 718
903
0718
0 1 1
500 1000 1500 2000
Yield Strength, MPa

FIGURE 20. Hydrogen E f f e c t s on Crack Growth in S t a i n l e s s Steels

- 51 -
TABLE 8

Stress Necessary for Slow-Crack Growth in Type 304L*

Stainless Steel

Net Section Time, Hours Crack

Stre!ss, MN/m^ Incremental Accumuljated Growth

600 325 325 No

641 72 397 No

682 72 469 No

724 72 541 No

765 72 613 No

786 1.4 (f-ailed) 614.4 Yes

* Crack developed during room temperature tensile test

in hydrogen environment; net section stress when
tensile test was stopped was 772 MN/m^. Specimen
then loaded in creep frame at indicated stresses
without removal from the hydrogen environment.

- 52 -
at threshold stress intensities of 110 to 44 MPa Vm7 The presence
of grain-boundary precipitates led to the lowest threshold stress
intensities. Slow-crack growth was also seen in Nitronic® 40,
where the threshold stress intensity was around 100 MPa VmT Sus-
tained load crack growth was not observed in HERF Type 304L, HERF
Type 316, or HERF Nitronic® 50 stainless steels at either 100 or
200 MPa hydrogen pressure.^^

Inconel® 718 is the only alloy for which hydrogen compatibilit

has been systematically evaluated by several test methods. " Me-
chanical tests included notched- and smooth-bar tensile tests, low-
and high-cycle fatigue, fracture mechanics, disc pressure tests,
and stress rupture. Ductility of notched tensile specimens and
fatigue life were more severely degraded than the other properties.
Notched tensile and disc pressure tests were sensitive to hydrogen
damage at pressures below about 6.9 MPa. These results and tests
on other materials suggest that notched tensile and disc pressure
tests are suitable for screening alloys for possible hydrogen ser-
vice. However, selection of a test method for more detailed in-
vestigation depends not only on sensitivity of the method and
severity of degradation but also on how well the test duplicates
loading conditions and hydrogen pressures for the intended service
application and possible material flaws.

Strain Rate. Hydrogen-assisted fracture in Type 304L

stainless steel displays an inverse strain rate effect. This
effect is common in all metals where hydrogen-assisted fracture
is observed. The more usual strain rate effect, observed in low-
temperature brittleness of carbon steels for example, is for higher
strain rates to increase brittleness. The inverse strain-rate
effect in hydrogen-assisted fracture has been cited as evidence
that hydrogen diffusion to high stress regions is necessary for
embrittlement.

FRACTOGRAPHY

Detailed examination of fractured specimens has been carried

out in recent years by scanning electron microscopy (SEM) which is
the basis for the descriptions of fracture modes in this section.

Fracture Modes

Ductile fracture of stainless steels occurs ordinarily by

a process of microvoid coalescence. Decohesion at the matrix-
inclusion interfaces or fracture of brittle inclusions creates
microvoids. The material between the voids fails by localized
flow or shear. This process originates throughout the individual
grains with no preferred fracture path, leading finally to the
common cup-cone or slant fracture.

- 53 -
 Hydrogen-assisted fracture of stainless steels is not associ-
ated with a unique fracture mode. On a macroscopic scale, frac-
tures are mixed mode, partly flat, and partly slant or a modified
cup-cone fracture. The microscopic fracture modes observed in
hydrogen-assisted fracture of stainless steels are microvoid coa-
lescence, cleavage, twin-boundary parting, intergranular separa-
tion, and interphase separation. In the following section, these
fracture modes are described and related where possible to alloy
composition, treatment, and external conditions, such as tempera-
ture and stress state.

Microvoid Coalescence

Most hydrogen-assisted fractures of stainless steels are, to

some extent, ductile, with areas of dimpled fracture evident even
when gross ductility is only 10 to 20%, Figure 21. With some al-
loys and test conditions, a change in dimple size is evident for
the HAF. Dimple diameters were measured for several Type 304L
stainless steel specimens, both with and without prior exposure to
hydrogen. The average dimple size in HAF was always smaller than
that for fracture of hydrogen-free specimens. Loss of tensile
ductility has been correlated with change in dimple size for aus-
tenitic steels. The larger the change in dimple size, the greater
the ductility loss. However, the results from tensile tests of
Type 304L steel at 250 and 270 K do not fall on the published
curve, which was derived from room temperature tensile tests.^'

Twin-Boundary Parting

Twin-boundary parting, in Figure 22, was the characteristic

mode of HAF in Type 304L stainless steel, Tenelon®, and Nitronic® 40
The incidence of this fracture mode was strongly dependent on tem-
perature and composition. For example, twin-boundary parting was
observed in Type 304L stainless steel but not in Type 316 stainless
steel. The most clearly defined examples of twin-boundary parting
occurred at test temperatures of 200 to 250 K in Type 304L stainless
steel, at 78 and 200 K in Tenelon® and at 200 K in Nitronic® 40.
Twin-boundary parting is observed in Tenelon® at 78 K even when
there is no hydrogen present.

Examination of numerous fractures has shown that twin-boundary

parting is characterized by the following features:

• A single facet extends over one grain only.

• Facets usually have steps about 5 pm in height.

• Traces of deformation bands in the underlying grain are visible.

- 54 -
 ••••• ; • jv- •-.,• * ^ 4 ^ ' :-^\

fe •••••^;V. ' ''••^•« •Strafe? ^'^

.Vi^

FIGORE 21, Dimpled Fracture of Type 304L Staiuless Steel

N^T
?P^lf'
/

0.» O, - • " • , ^ > ^ ^'-.^

."-.

4if %SS
/'?
1 •'<*

FICIIIE 22, Twia-Bonsdarj fsrSiag ia Type 3041. Stainless Steel.

High Energy Eat® Forged, Tested at 200 K.

- 55
m Opposing halves of the fracture match and interlock.

® Facets are slightly curved because of lattice bending prior to

separation along the twin boundary,

m River patterns characteristic of cleavage are never seen.

On a longitudinal section of a tensile specimen, twin-boundary

parting shows as microcracks throughout the deformed portion of the
specimen. Microcracks form on boundaries oriented roughly trans-
verse to the tensile axis, Figure 23. The frequency of occurrence
of microcracks correlates with the temperature of maximum hydrogen
damage, Table 9, just as the number of facets on the fracture face
varies with temperature. Figure 24. Multiple nucleation of micro-
cracks is seen also. Figure 25, and suggests that twin-boundary
parting proceeds by the linking of many cracks which nucleate at
intersections of deformation bands with the twin-boundary. Acous-
tic emission confirms the notion that crack growth is by many small
steps rather than one abrupt energy release.

Transgranular Cleavage

Transgranular cleavage was a form of HAF observed in ferritic

and martensitic stainless steels. For example, the fracture mode
of 17-4PH was predominatly transgranular cleavage as seen in
Figure 26. This fracture mode can yield facets similar to those
of twin-boundary parting, but the facets are readily distinguished
at high magnification by the occurrence of river patterns.

Intergranular Separation

Intergranular separation was characteristic of HAF of alloys

containing more than about 30 or 35% nickel but was not limited to
these alloys. Commercial grades of nickel displayed the most
clearly delineated intergranular separation. Figure 27. However,
specific thermal treatment could evoke intergranular separation in
most alloys. In the case of nickel, intergranular separation was
increased by prolonged heating at 770 K.

Intergranular separation has been observed in C-shaped speci-

mens of HERF Nitronic® 40 and A-286 stainless steel also. Figure 28.
Precharging with hydrogen is needed to evoke the intergranular
separation. The presence of intergranular fracture in C-shaped
specimens and its absence in smooth bar tensile specimens suggests
that stress state affects fracture mode in stainless steels. There
is a bending moment on the C-shaped specimens which is absent in a
tensile test.

- 56 -
 •

'- I

FIGIJIE 23, Microcracks M o n g Boundaries ©f Annealing Twins

•
- 57 -
TABLE 9

Occurrence of Microcracks in Hydrogen-Saturated

Austenitic Stainless Steel

Test Microcrack Frequency (Number Cracks,

Temp,_K Type 304L Tenelon® Nitronic® 40

380 0 _ 0
348 - 0 -
298 - - 60
273 90 900 _

248 150 _ 230

198 1800 1100 1400
78 0 (4)* 0

* Number of microcracks observed in close proximity to

the fracture surface.
All specimens saturated with deuterium at 69 MPa
pressure at 570-620 K.

- 58 -
.r
198 K 273 K 2« K
FIOTIE 24. fariatioa @f F®€®t Appearaac® with T@®t
Temperature in Type 304L Stainless Steel

- 59 -
 * -r ^ »
,^.-«

1 ^ -"t
f
*>

a) On Twin Boundary

*;».•'• • -•. ••

•-&?•.• •-'

%-'\v.-\.-

* • * . . . , -

" .' »i

b) On Grain Boundary

FIGUSE 25. Multiple Crack Mucleation Along lotrndsries

- 60 -
 ^m-

FXGUIE 26. Transgramslar €le@ir@§@ in l?*^ n

Frecipitation-Hardenable Stainbles* Steel

•«•**

i
T
I
{

C.r.

flGUlE 27. Intergranular 8«p®r®ti@® in Mickel and Inconel® 718

- 61 -
 m*

,(jj

FICTK 18. B I F •itr@®ie» 40, Orieetatioa 1.

Integranular Fracture in Hydrogen,

- 62 -
 Intergranular fracture paths were observed after tensile tests
in a hydrogen environment in sensitized specimens of Type 304Lj 304,
and Nitronic® 40 stainless steel but not Type 309S. ^ Moreover, the
presence of an intergranular fracture does not appear to be contin-
gent upon apparent continuity of the carbide phase. Continuous
carbides and intergranular fracture were observed in Type 304 and
Nitronic® 40 stainless steels, but no intergranular fracture occur-
red in Type 309S stainless steel in spite of formation of continu-
ous grain-boundary carbides. Furthermore, extensive intergranular
fracture took place during failure of Type 304 stainless steel
where carbide precipitation was discontinuous. In the case of
Type 304L steel, a distinct difference in fracture appearance was
noted which depends upon the frequency of occurrence of grain
boundary carbides. An intergranular fracture path occurred only
when the carbide network was nearly continuous on at least some
grain faces. Otherwise, the fracture was mixed void coalescence
and partially brittle but not intergranular.

Although continuity of carbide precipitation on grain faces

appears to be the dominant factor controlling the fracture path,
one or more other factors must be present to explain the above
observations. Phosphorus segregation to the grain boundaries
and possible formation of an (Fe, Cr) Ni-P surface phase have
been demonstrated as the cause of temper embrittlement of Type 304
stainless steel.-^" Intergranular fracture at 150 K was observed on
notched tensile specimens annealed 2 hours at 920 K. Surface seg-
regation of phosphorus is also found by Auger electron spectroscopy
after vacuum annealing at 820 to 1020 K. These studies suggest
that phosphorus may be a factor in causing intergranular fracture
of sensitized austenitic steels, but experimental data are inade-
quate to confirm the possibility.

Interphase Separation

HAF in austenitic steels often propagates along interphase

interfaces such as austenite-martensite or austenite-ferrite inter-
faces. Welds in austenitic steels, such as Type 304L, contain
several percent of S-ferrite which forms during solidification of
the weld bead. The presence of the 6-ferrite helps prevent hot
shortness or weld cracking but is a preferred path for crack propa-
gation in the presence of hydrogen. Figure 29a.

Some stainless steels are unstable when deformed and transform

to e-martensite (a HOP lattice) or a-martensite (a BCC lattice).
Again the interfaces between the austenite and either form of
strain-induced martensite are likely crack paths. Fracture surfaces
in these cases have a crystallographic appearance with many inter-
secting planar facets. Figure 29b. In Type 304L steel, this frac-
ture path is common around room temperature in specimens that have
been precharged with hydrogen prior to deformation.

- 63 -
) Austenite-Ferrite Interface in Type 304L fcM ^ t a l

b) Austenite-Martensite Interface in Type 304L

FIGURE 29. Fracture Along Interphase Interfaces

in Austenitic Steels

- 64 -
 Fracture along austenite-martensite interfaces in Type 304L
stainless steel appears more common around room temperature than at
lower temperatures. Twin-boundary parting dominates the fracture at
200-250 K and microvoid coalescence is seen at 78 K, Clearly, the
operative fracture path is very sensitive to test temperature.

Correlation of Fracture Mode With Composition

Fracture modes for HAF of iron-chromium-nickel alloys vary

with percent of nickel in the alloy. Figure 30. Alloys such as
17-4PH with only 4% nickel fail by transgranular cleavage. As the
nickel content increases, the steels become austenitic (metastable),
and fracture by twin-boundary parting appears at around 10% nickel.
A region of dimpled fracture occurs from 15 to 25 or 30% nickel,
which is gradually supplanted by intergranular separation as the
nickel content increases. However, there were only a few alloys in
which only a single fracture mode was observed. In most instances,
two or more fracture modes occurred, one of them being microvoid
coalescence.

Alloy composition is a controlling factor in HAF in two ways:

1) the base alloy determines slip character and phase stability
during straining, and 2) impurity and trace elements may be strongly
segregated and induce intergranular separation either alone or in
combination with hydrogen. Planar slip is associated with low-
nickel austenites and leads to high-stress concentrations at slip
barriers, such as inclusion, twin, and grain boundaries. Sites of
high-stress concentration may act as microcrack nucleation centers,
especially in the presence of hydrogen, which may lower the cohe-
sive strength, Austenite stability under strain is also correlated
with nickel content. Formation of deformation twins, e-phase, and
a-martensite occurs more readily with lower nickel concentrations.
Both slip planarity and phase stability are related to stacking-
fault energy. Greater resistance to hydrogen damage in austenitic
steels has been correlated with higher stacking-fault energy. The
transition from twin-boundary parting to dimpled fracture at around
12 to 14% nickel correlated with an increase in stacking-fault
energy to over 30 or 40 mj/m^.

Intergranular separation in all of the stainless steels may be

attributable to impurity segregation. Sulfur in high-nickel alloys
and phosphorus in austentic steels are known to cause intergranular
failures. These impurities are not the only causes of intergran-
ular fracture, however. Sensitization of austenitic steels leads
to carbide pecipitation and grain-boundary regions depleted in
chromium and causes intergranular fracture in Type 304L stainless
steels. These examples illustrate the sensitivity of HAF to rela-
tively small changes in either base alloy composition or impurity
content. Control of melting, casting, and mechanical and thermal

- 65 -
 Fracture Mode
Fe (D) Dimple

( T ) Twin-Boundary Parting

(C) Cleavage

(T) Intergranular

Ni 0:'
^yy>^mmmmmmm.
/,-j(i^-_.;-,._-,-/_-•,-•,•••.,• •^.;••••• ••••/• ••••• v . . , . . . . / . . . y . \ . . / . \ . y ; . ^ y

FIGURE 30. Composition Regimes for Fracture Modes Observed in

HAF of Iron-Ghromium-Nickel Alloys

- 66 -
processing becomes more important for hydrogen service than for
service in air. Small changes in local composition due to varia-
tions in process control can develop conditions for HAF in an
otherwise resistant alloy.

HELIUM EMBRITTLEMENT

The hydrogen isotopes are alike in their short-term chemical

effect on the properties of stainless steels. Long-term effects
differ, however, because tritium decays with a 12.35 year half-life
to helium-3 by emission of a low energy (5.7 keV average) beta
particle, whereas protium and deuterium are stable. Helium is
known to embrittle many metals but normally elevated temperatures
are required for an observable effect.^' Exposure of metals to
tritium causes "hydrogen damage," which occurs immediately and
slowly diminishes as tritium decays, as well as helium embrittle-
ment, which becomes significant after tritium decay has generated
sufficient helium. Helium damage is additive to the damage caused
by the hydrogen isotopes and has been identified in Armco® (Armco,
Inc.) iron - and niobium as well as in several grades of stainless
austenitic steel.

The earliest studies of the effects of helium-3 on the mechan-

ical properties of stainless steel were reported in 1975.^'»^°
Foil specimens (0.025 cm thick) were exposed to 47.5 MPa tritium
at 343 K for 17 months; helium build-in was estimated at 6.2 mol
He(STP)/m^. Subsequent tests at room temperature indicated that
the exposure increased the yield strength, but had little effect on
the ductility of the Type 3098 steel specimens, although the
ductility of Type 304L steel was lowered. Microvoid coalescence
was the primary fracture mode. Helium bubbles about 5 nm in size
were found in both steels in grains and on grain boundaries.

Transmission electron microscopy of samples charged with trit-

ium showed gas bubbles formed on dislocation networks (Figure 31),
and grain boundaries following a 0.5 hour anneal at 973 K. Further
a tensile test at 973 K showed that the helium, which had built in
during the charging, agglomerated during the test, which caused the
fracture mode to change from ductile rupture to intergranular sepa-
ration and reduced the strength and ductility of both Type 3098 and
304L stainless steels.

Helium damage at ambient temperature was convincingly demon-

strated in "smooth-bar tensile specimens (2.5-cm gage length and
0.356-cm diameter) of HERF Nitronic® 40 stainless steel gas-phase
charged with tritium at 64 MPa at 470 K for 1450 hours and aged

- 67 -
FIGURE 31. Helium Bubbles in Type 304L Stainless Steel.
Tritium Charged, Aged, and Annealed at 1273 K

- 68 -
5.5 years at 250 K. Comparison specimens of Type 304L stainless
steel were only affected a small amount by the same treatement.^'
Tritium and helium distributions in the specimens are nonuniform
for these charging conditions. Some of the charged specimens were
tested at 298 K immediately after charging, and others were stored
from five to seven years at 250 K before testing, Tables 10 and 11.
After 5.5 years' storage, helium concentrations were about 100 mol
He(STP)/m^ at the surface of the Type 304L steel specimen and
about 174 mol He(STP)/m^ in the Nitronic® 40. Average concentra-
tions over the specimen cross section were 31 and 52 mol He(STP)/m3

These test results show:

1) Hydrogen and tritium reduce ductility of both Type 304L and

Nitronic® 40 stainless steel (Tables 10 and 11).

2) There was only a small change in the mechanical properties and

no change of fracture mode in Type 304L stainless steel with
100 mol He(STP)/m3.

3) A drastic reduction in ductility and a change in fracture mode

from microvoid coalescence to intergranular separation occurred
in Nitronic® 40 stainless steel (Table 11) with 174 mol
He(STP)/m3.

There was a high density of small defects in the tritium-

charged and aged Nitronic® 40 stainless steel, Figure 32. Contrast
analysis of the strain fields produced by these defects indicated
that they were helium bubbles approximately 5 nm in size. The size
and density of bubbles (2 x 10^" bubbles/cm^) was consistent with
the estimated helium content of the specimen. Measurements of the
images produced by these defects indicate a misfit strain of about
0.017. Rather than being located at grain boundaries, the bubbles
were dispersed throughout the matrix. A few bubbles were found in
the Type 304L stainless steel specimens, in keeping with the lower
helium content. Helium is known to cause intergranular fracture by
the formation of intergranular helium bubbles at high temperatures
(T/Tm = 0.5), where helium migration to internal boundaries can
occur.^' Data for Types 304L and 3098 stainless steel and
Armco® (Armco, Inc.)^" iron charged with tritium have shown
helium embrittlement with only 0.2 mol He(STP)/m^ can occur after
a high-temperature anneal or testing at a high temperature.

HERF Nitronic® 50 and welds in Type 304L stainless steel have

shown ductility losses associated with helium-3.^-^ Calculated
helium concentration profiles in these specimens were similar to
those in the HERF Nitronic® 40 and HERF Type 304L stainless steels
but the surface concentration of tritium was higher. All fractures
in the HERF Nitronic® 50 and welds of Type 304L stainless steels
were ductile.

- 69 -
TABLE 10

Mechanical Properties of HERF Type 304L Stainless Steel

Total Plastic
Exposure Strengt h, MPa Elongation, Strain to
Conditions Yield Ultimate % Fracture

None 630 790 32 1.83

69 MPa H2 at 480 K, 600 840 35 0.87

55 days

47 MPa T2 at 480 K, 620 745 33 0.73

17 months and aged
5.5 years at 250 K

TABLE 11

Mechanical Properties of HERF Nitronic* 40 Stainless Steel

Plastic
Exposure Strength, MPa Total Strain to
Conditions Yield Ultimate Elogation, % Fracture

None 800 930 29 1.20

30 MPa H2 at 480 K, 845 990 28 0.76

56 days and aged
7 years at 250K

65 MPa T2 at 480 K, 830 940 32 0.73

60 days

47 MPa T2 at 344 K, 870 925 0.10

17 months and aged
5.5 years at 250 K

- 70 -
 ORNL - Photo 3598-80

a) Hydrogen charged 6 9 M P a at 470K b) Tritium charged 47.5 MPa at 345K

and aged 5.5 years and aged 5.5 years

FIGURE 32. Microstructures i n HERF N i t r o n i c * 40

- 71 -
 A new investigation of helium effects was begun and is still
in progress.^" Notched C-shape specimens 3.8 mm thick and 25 mm
outer radius were machined from five stainless steels: Type 304L,
Type 316, Nitronic® 40, A-286, and a modified A-286. The modified
A-286 was in the annealed condition, all others were HERF. Speci-
mens were not fatigue precracked. The specimens were exposed to
tritium at 61 MPa pressure at 422 K for six months.

Calculated tritium concentrations in Type 304L stainless steel

were 270 mol T2(STP)/m3 at the surface and 60 mol T2(8TP)/m3
at the center based on the hydrogen solubility equation of Louthan
and Derrick.^-' Tritium and helium contents average over the cross
section were 167 mol T2(8TP)/m^ and 10 mol He(8TP)/m^. Following
the 15-month ageing, average tritium and helium contents were 160 mo
T2(STP)/m3 and 24 mol He(STP)/m3. Actual concentrations in
Nitronic® 40 are probably higher since measurements of hydrogen
solubility at 470 K have shown that hydrogen solubility is 70%
higher than in Type 304L steel.^^

Tensile tests were run immediately after exposure and follow-

ing 15 months' storage at 273 K. Control specimens were exposed to
air at 422 K for six months and stored at low temperature. Tensile
tests were made in air at room temperature and the load-deflection
curves were analyzed by the J-integral technique.^^ J-integral
at maximum load (Jm) and deflection at maximum load are given in
Table 12 for the several test conditions.

The data demonstrate that helium has degraded the mechanical

properties of all five steels. Immediately after the six-month
exposure, the difference in Jm between the control and charged
specimens was due mostly to the tritium with only a small helium
effect. After 15 months' storage at 273 K, the difference between
charged and control specimens was due to the additional helium gen-
erated during storage. HERF Type 316 and HERF Type 304L stainless
steels were affected the least by the helium, and HERF Nitronic® 40
and annealed A-286 (modified) were affected the most. The occur-
rence of a large helium effect in Nitronic® 40 was not surprising,
as a comparable result was observed earlier.^" The difference
in response to the presence of helium shown by HERF A-286 and
annealed A-286 (modified) appears to be associated with microstruc-
tural differences arising from prior thermal or thermo- mechanical
treatments. Composition differences between the two varieties of
A-286 are small and probably contributed little to the relative
helium effect.

Sustained load tests were initiated after 22 months storage on

all steels except modified A-286. Loads corresponded to 80 percent
of the fracture load that was measured after 15 month's storage.
Rapid crack growth occurred in Nitronic® 40 and A-286 stainless
steels but has not yet been detected in Type 316 or Type 304L after
five months under load. The latter tests are continuing.

- 72 -
TABLE 12

Tritium and Helium Effects on Fracture Toughness of Stainless Steel

Fracture
Deflection at Toughness
Alloy History* Max. Load, mm Jm, kJ/m2

304L C 3.71 1280

HERF T 1.55 360
C + A 3.33 1120
T + A 1.32 190

A-286 C 1.95** 610**

HERF T 1.80 530
C + A 1.70 350
T + A 1.14 130

Nitronic® 40 C 2.67 1200

HERF T 1.50 480
C + A 2.62 960
T + A 1.02 75

316 C 3.51** 1500**

HERF T 2.39 890
C + A 2.97 950
T + A 2.03 630

A-286 C 2.86 840

(modified) T 1.50 350
Annealed C + A 2.79** 680**
T + A 0.84** 70**

* C - Control
C + A - Control + Aged
T - Tritium Charged
T + A = Tritium Charged + Aged.
** Single Specimen. All others in duplicate,

- 73
 The fracture mode tended to change from void coalescence to
intergranular in all alloys except HERF Type 316. Fracture of
annealed A-286 (modified) was partly intergranular after tritium
charging and entirely intergranular after subsequent aging. Inter-
granular fracture was evident in isolated areas of the Nitronic® 40
after tritium charging and became dominant after aging. Both
Type 304L and Type A-286 stainless steels failed by mixed fracture
(void coalescence and intergranular) after charging and aging,
whereas only void coalescence was observed in the tritium charged
specimens of these alloys.

Helium and hydrogen (tritium) damage occurred simultaneously,

and no attempt was made to factor the measured damage between the
two causes. The magnitudes of the two effects will change with
time, as the tritium decays to helium. The data accumulated so far
suggest however, that helium damage is more potent than hydrogen
(tritium) damage for equal atomic concentrations of the two ele-
ments . There were decreases in Jm of 50 to 90% as a consequence of
the 15-month aging during which time the average helium content
increased by a factor of 2 and the average tritium content de-
creased by 4%.

Distribution of the helium within the specimens has not been

established as yet. In the Nitronic®-40, helium is presumably dis-
tributed in fine bubbles throughout the matrix as in the earlier
specimens. There is no reason to believe that this same helium
distribution should apply to the other alloys, particularly when
their widely differing microstructures are taken into account. For
example, fine carbo-nitride particles occur more commonly in
Nitronic® 40 than in Type 304L or Type 316 stainless steels, and
both varieties of A-286 may contain gamma prime (Ni3[Al, Ti], eta
(Ni3Ti) and TiC phases in varying quantities and distributions.
All of these phases could act as traps for tritium and thus create
localized high helium concentrations or bubbles which would be dis-
tributed within the alloy in the same manner as the phases. Sever-
ity of helium damage would be related then to trap characteristics
such as, quantity of trapped tritium at the interface, interface
coherency and whether the trap were reversible or irreversible.
However, the relation between helium damage and the trap character-
istics need not be the same as between hydrogen damage and trap
characteristics because the strain fields of helium and hydrogen
differ and hydrogen is chemically reactive.

- 74 -
CONCLUSIONS

Substantial progress has been made in identifying important

material and processing variables which influence hydrogen compat-
ibility of stainless steels. The results provide guidance in alloy
selection and thermomechanical processing for a limited range of
hydrogen pressures (to 69 MPa), operating temperatures (78-400 K ) ,
and mechanical loading conditions (constant or rising load). How-
ever, there are service conditions for which data are inadequate to
allow selection of a reliable alloy or prediction of its service
life.

Several unresolved issues remain that are important to util-

ization of stainless steel in environments where hydrogen isotopes
are present:

• The relative importance of base alloy composition, precipitate

structure and morphology, and mechanical processing in attaining
high strength (>900 MPa yield strength) with good hydrogen com-
patibility (Kfjj >80 MPaVm) is not certain.

• Selection of alloys for low hydrogen pressure (O.OI to 10 MPa)

applications has not been systematically examined. Pressure
limits and acceptable temperature ranges have not been corre-
lated with service life and stress states for low pressure
service, particularly at elevated temperatures (T >400 K) where
hydrogen attacks occur and creep conditions are encountered.

• Segregation of alloy elements, impurities, and hydrogen can

adversely affect hydrogen compatibility in an alloy otherwise
acceptable for a given service. The relationships of segrega-
tion to hydrogen compatibility on the one hand and to processing
variables and control on the other have not been examined in
detail.

• Mechanical test techniques need to be analyzed and compared to

permit selection of those methods which best relate to and pre-
dict hydrogen compatibility for given service conditions (load-
ing, hydrogen pressure, and temperature). J-integral methods
were investigated to characterize ductile fracture and hydrogen
degradation of tough alloys such as austentic stainless steels.
Initial results were encouraging and suggest continued develop-
ment of experimental and analytical procedures.

- 75 -
• Mechanisms of hydrogen-assisted fracture and low temperature
helium embrittlement are not known. Several hypotheses have
been advanced to account for each form of degradation, but there
are as yet no generally accepted explanations for these
phenomena.

• Material, processing, and environmental conditions controlling

helium embrittlement are not well known. Experiments have begun
to examine systematically the phenomenon of low temperature
helium embrittlement.

Investigation of hydrogen effects on the mechanical behavior

of stainless steels is continuing at the Savannah River Laboratory.
The issues identified above serve in part to guide the research
activity.

- 76 -
ACKNOWLEDGEMENTS

This report includes work by several investigators in the Savannah

River Laboratory whose contributions are acknowledged: T. L. Capaletti,
M. R. Dietrich, J. A. Donovan, C. W. Krapp, M. R. Louthan, Jr.,
D. A. Mezzanotte, L. P. McCabe, D. E. Rawl, Jr., and P. E. Zapp. The
high-purity stainless steels were supplied by N. F. Fiore, Notre Dame
University. Transmission microscopy of Nitronic® 40 stainless steel,
Figure 32, was performed by P. S. Sklad, Oak Ridge National Laboratory.

- 77 -
ALLOY DATA SHEETS

ALLOY INDEX

Nominal compositions of alloys are shown in Table A-1. Nomina

tensile properties are shown in Table A-2. Measured properties
under various experimental conditions are summarized in the data
sheets listed below.

I. Iron-Chromium-Nickel Alloys Data Sheet

304L IA-1 to IA-16

304N IB-1
309S IC-1
310 ID-1 to ID-2
316 IE-1
Carpenter 20 Cb-3 IF-1
Incoloy® 800H (Huntington Alloys, Inc) IG-1
Nickel 200 IH-1 to IH-2
Nickel 301 IJ-1 to IJ-2
440 C IK-1

II. Iron-Chromium-Nickel-Manganese Alloys

Tenelon® (U.S. Steel Corp) IIA-1 to IIA-3

Nitronic®-40 (21-6-9)(Armco, Inc) IIB-1 to IIB-12
Nitronic®-50 (22-13-5)(Armco, Inc) IIC-1 to IIC-3
18-18 Plus® (Carpenter Technology) IID-1
X18-3 Mn IIE-I
18-2 Mn IIF-1
216 I IG-1

III. Precipitation Hardenable Alloys

A-28 6 IIIA-1 to IIIA-5

JBK-75 IIIB-1 to IIIB-2
17-4PH IIIC-1 to IIIC-2
AM-350 IIID-1
AM-363 IIIE-1
CG-27 IIIF-1
Ni-SPAN-C (Alloy 902) IIIG-1

IV. High Purity Alloys

A - 18Cr-10Ni IVA-1
B - 18Cr-14Ni IVB-1
C - 18Cr-19Ni IVC-1
- 78 -
TABLE A-1

Nominal Alloy Composition, wt %

Alloy Cr Ni Mn Mo Other

F e - C r - N i - Alloys
304 L 19 10 - -
304N 19 9 - - 0.13 N
309S 23 13 - -
310 25 20 - - 0.25 C
316 17 12 - 2.5
440C 19 - - 0.75 0 . 9 5 t o 1.20 C
C a r p e n t e r 20 (Cb-3 20 34 - 2.5 3 . 5 Cu, 06 Nb
I800H 21 32 - - 0 . 7 5 Cu, 0 . 3 A l , 0 . 3 T i
1718 19 52 - 13 5 (Nb + T a ) , 1 T i , 0 . 5 A
Ni200 - 99 + - -
Ni301 - bal - - 1 S i , 4 . 5 A l , 0 . 6 Ti

F e - C r - N i - 1to-N Alloy s
216 20 6 8 2 0.32 N
Tenelon® 18 - 15 -
N i t r o n i c ® -40 21 6 9 - 0.15 to 0.4 N
N i t r o n i c ® -50 22 13 5 2 0.2 to 0.4 N
18-18 Plus® 18 0.5 18 1 0 . 4 N, 1 Cu, 0 . 1 Co
X 1 8 - 3 Mn 18 3 12 - 0.3 N
1 8 - 2 Mn 18 2 13 -

P r e c i p i t a t i o n --Hardenable Alloys <

1 7 - 4 PH 16.5 4 - - 4 Cu, 0 . 3 Nb
A-286 15 26 - 1.25 2 Ti, 0.25 Al, 0.3 V
JBK-75 15 30 — 1.25 2 T i , 0.25 Al,
0 . 0 0 1 B, 0 . 2 5 V
AM363 11.5 4.5 - - 0.5 Ti
CG27 13 38 - 6 2 . 5 T i , 1.6 A l , 0 . 6 Nb
AM350 16.5 4.3 - 2.8 0.1 N
Ni-SPAN-C
A l l o y 902 5 42 - - 0 . 5 A l , 2.5 Ti

H i g h - P u r i l ty A l l o y s
A 18 10 - - N <0.01 in a l l
B 18 14 - - three alloys
C 18 19 - -

- 79 -
TABLE A-2

Nominal Tensile Properties (Annealed Material

Unless Otherwise Noted)

Strength, MPa
Alloy Yield* Tensile Elongation, %

304L 230-•270 540-560 55-60

304N 290-•330 620 50-55
309S 275-•310 620-650 45
310 310 650 45-50
316 207-•290 550-585 45-50
440C 450-•1890 760-1965 2-14
Carpenter 20 Cb-3 250 600 50
I800H 140-•345 450-650 30-50
1718 118C)-1250 1350-1400 16
Ni 200 103-•207 380-550 40-55
Ni 301 210-•1200 620- 1450 15-55
216 428. 745 50
Tenelon® 570 930 56
Nitronic®--40 414 690 40
Nitronic®-•50 448 828 45
18-18 Plus® 520 900 60
X 18-3 Mn 580 810 45
18-2 Mn 730 1000 51
17-4 PH 940 980 5
A-286 760 1100 25
JBK-75** 800 1090 14
AM 350 420 1160 70
AM 363 890 890 7
CG 27 810 1160 29
Ni-SPAN-C 760-•870 900- 1200 6-25
Alloy 902

* 0.2% offset.
** HERF & Age.

- 80 -
IRON-CHROMIUM-NICKEL ALLOYS

DATA SHEET IA-1

Type 304L Stainless Steel Bar Stock, As Received*

Test Condition Hydrogen** Strength, MPa Elongation, % Fractu

Temp, K Environ. Exposure Yield Ultimate Uniform Total Strain

380 AIR NONE 240 680 58 69 1.78

69 MPa 260 730 60 70 1.27

273 AIR NONE 310 1160 80 89 1.56

69 MPa 330 870 44 44 0.45

200 AIR NONE 360 1500 61 70 1.27

69 MPa 390 1210 44 44 0.25

78 LN NONE 390 2200 60 64 1.27

69 MPa 430 2100 59 65 1.27

* Heat Analysis, Appendix D-1; Tensile B, Appendix C-2.

** Exposure conditions: 69 MPa at 470 K for 1449 days.

DATA SHEET IA-2

Type 304L Stainless Steel, As Received

Test Conditions Hydrogen Impact

Temp, K Environ. Exposure Energy, J

298 AIR NONE 194

17.9 MPa* 185

78 AIR NONE 165

17.9 MPa* 110

* 17.9 MPa hydrogen pressure at 470 K for 1000 hours.

- 81 -
DATA SHEET IA-3

Effect of Test Environment on Tensile Properties

of Type 304L Stainless Steel Tubes*

Exposure Conditions Tensile Properties

Gas Temp, K da

He 425 32 270 560 59

H2 425 32 320 480 19
T2 425 32 300 490 22
H2 425 8 260 490 26
T2 425 8 250 490 22

* All tensile tubes tested at room temperature with 69 MPa gas;

data reported are averages of at least two samples.

DATA SHEET IA-4

Tensile Properties of Type 304L Stainless Steels Containing

Hydrogen and Helium

Test Condition Hydrogen Streng th, MPa Elong

Temp, K Environment Exposure Yield Ultimate %

300 Air none 327 734 49-56

300* Air ** 400 733 28-32
300* Air ** 434 744 28
973t Air none 152 237 31
973t Air ** 179 190 1.5

* Specimens contained tritium and Helium-3.

** 328 mol hydrogen isotopes and 6.2 mol helium per m^ metal.
t 146 mol hydrogen isotopes and 25 mol helium per m^ metal.
Held 1/2 hour at 973 K before testing.

- 82 -
DATA SHEET IA-5

Type 304L Stainless Steel, High Energy Rate Forged*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environ. Exposure** Yield Ultimate Uniform Total Strain

380 Air None 440 630 32 44 1.72

69 MPa 440 650 32 43 1.63

298 Air None 480 930 57 68 2.00

69 MPa 510 990 55 62 0.95

250 Air None 490 1100 52 61 1.65

69 MPa 610 1120 41 41 0.40

200 Air None 660 1390 46 55 1.37

69 MPa 620 1300 43 44 0,38

* Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks.

DATA SHEET IA-6

Type 304L Stainless Steel, High Energy Rate Forged*

Test Condition Hydrogen Impact

Temp, K Environ. Exposure Energy, J

298 Air None 199

298 Air 29.6 MPa H2** 152

77 Air None 160

77 Air 29.6 MPa H2** 95

* Impact, Appendix C-8.

** Exposure of 56 days at 470 K.

83 -
 h4 a
1 >
kJ/m^ S) H
o n
n 1 a' o >
C 09
rr ET -p. en CXI o e so
fB 0> 1
'^ "O
o o o o o o n
H- (B
o o o o o o H
c a. PI
rr
O (B
3- 3
P> CO n nn i ?
"1 H - O -D Duplicate Tests Agreed Exactly rr
00 I-- -J ai
_ j . n
(D (B -s
W (B cu ^5 rD ro 1
01 13 —*
r+ o; CO l/i
rt >
r* — <
01 n> rB T r+ <-<•
•V r+ — I -5 l O fB fD o
OVX) —1. -••CO Q. O. 1
O 3: C Q.
vo m H
3 3 '< ? Z 3 3
a.
-s a:
^P-. o '< (B
to X vo Q. —•
fB -5 ->•
a> d 3 O C
1 rt 1 VO 3 o
>j fB
00 <3v • -S <J1
f^
•p- NJ One Specimen Tested
O fB CO
1 H 3
P«< f t n
CO 0)
Hi rr H'
O O 9
H 1— 3 >-'
3
u> n
Ov
« VO «
fB CO
fB S rr
?r -TS A
CO 01
A
o vo Only One Specimen Tested
ES -5 O
fB (B H»
3 OQ
O SP
1
EC H
N> O
3
n
OQ
*<

Id
o
1
OQ

n
»
DATA SHEET IA-8

Fracture Parameters for Type 304L Stainless Steel*

Tested In Helium

V "" 1 Tested in Hydrogen

1 Charged with Deuterium, Tested in Hydrogen

400 - -

f^ • •
300 -

^'h
//
/ %/ .
//y -i-l
••
-o *
V/ VA

'1' 1 i
200
-o X!
OJ
+-'
(LI 0;
H- 1-
f
'
c
U
OJ
u
D.
100 - oo
OJ
OJ OJ
c c
o o
>, ' >,
c c
o o

V/ ;v:
"arallel 45°' 90°
Orientation Orientation Orientation

* C-shaped t e n s i l e . Appendix C-7. Test i n 69 MPa He or H2.

Deuterium charge a t 69 MPa a t 620 K for 3 weeks.

- 85 -
DATA SHEET IA-9

Effect of Heat Treatment on Mechanical Properties of

Type 304L Stainless Steel*

Heat Test Strength, MPa Elongation, % Fracture

Treatment Environment Yield Ultimate Uniform Total Strain

As-received 69 MPa He 390 930 62 71 2.21

GS = 9.5 ym 69 MPa H2 390 910 56 62 1.43
1170 K-24 hrs 69 MPa He 260 970 82 89 2.30
GS = 30 pm 69 MPa H2 240 970 88 94 1.71
1270 K-24 hrs 69 MPa He 250 970 90 99 2.30
GS = 55 \im. 69 MPa H2 240 930 86 91 1.17
1470 K-24 hrs 69 MPa He 190 830 81 88 2.21
GS = 340 pm 69 MPa H2 180 830 84 88 1.05

* Heat analysis; Appendix D-1; Tensile B, Appendix C-2.

DATA SHEET IA-10

Grain Size Dependence of Mechanical Properties - Test at 220 K*

Hydrogen Grain Strength, MPa Elongation, % Fracture

Exposure Size, pm Yield Ultimate Unit Total Strain

None 6.4 520 1310 56 63 1.71

42.0 340 1210 .60 72 1.70
290 250 1130 55 63 1.64

69 MPa** 6.1 630 1040 35 35 0.27

26 400 1020 47 47 1.10
50 370 860 37 37 0.40
260 270 690 31 31 0.39

* Heat analysis. Appendix D-1; Tensile B, Appendix C-2.

** Average deuterium contents measured on samples from
the tensile specimens were 4.7 ccD2/cc (69 MPa).

- 86 -
DATA SHEET IA-11

Effect of Deformation Rate on Hydrogen Damage*

TYPE 304L Stainless Steel

T = 220 K
Grain Size: 6 m

Cross
Hydrogen Headspeed, Strength, MPa Elongation, % Fracture
Exposure** mm/min Yield Ultimate Uniform Total Strain

None 51 570 1170 50 60 1.68

0.51 5.20 1310 56 63 1.70

69 MPa 51 675 1210 52 52 0.92

0.51 630 1040 35 35 0.27

* Heat analysis. Appendix D-1; Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks.

DATA SHEET IA-12

Mechanical Properties of Sensitized Type 304L Stainless Steel*

(Smooth Bar Tensile Specimens)

Test Strength, MPa Fracture

Treatment Environment** Yield Ultimate Strain

Solution Air 380 630 2.00

Anneal Helium 375 600 2.20
Hydrogen 370 580 1.38

Sensitized Air 300 560 1.78

Helium 350 670 1.90
Hydrogent 330 660 0.70
Hydrogentt 350 660 0.80

* Heat analysis, Appendix D-1; Tensile B, Appendix C-2.

** 69 MPa gas pressure.
t Nearly continuous carbide network on some grain boundaries,
tt Isolated carbides.

- 87
DATA SHEET IA-13

Mechanical Properties of Notch Bar Tensile Specimens of

Type 304L Stainless Steel*

Test Streng th, MPa Fracture

Treatment Environment** Yield Ultimate Strain

Solution Airt 700 750 0.41

Anneal

Sensitized Airtt 350 590 1.26

Heliumtt 410 740 1.10
Hydrogentt,1 430 590 0.35
Hydrogen 11 480 620 0.38
Airt 510 680 1.17
Heliumt 540 790 1.00
Hydrogent,1 730 750 0.30
Hydrogen 11 690 0.20

* Heat analysis. Appendix D-1 and Appendix C-2.

** He and H2 at 69 MPa. Air at 0.1 MPa.
t Deep notch.
tt Shallow notch.
1 Nearly continuous carbide network on some grain boundaries,
11 Isolated carbides.

- 88 -
DATA SHEET IA-14

Effect of Hydrogen Charging on Notch Bar Tensile Properties

of Type 304L Stainless Steel*

Nominal TIensile Fracture

Condition Specimen Strength. MPa Strain

As received Smooth 600 1.50

Notch 770 0.30

Annealed** Smooth 600 1.43

Notch 710 0.24

Hydrogen chargedt Smooth 530 0.37

Notch 580 0.13

* Tensile C, Appendix C-4.

** Annealed 200 days at 380 K in argon,
t Exposed to hydrogen gas at 69 MPa for 200 days at 380 K.

DATA SHEET IA-15

Type 304L Stainless Steel Notch Tensile Strength*

Notch Tensile
Test Strength,
Environment MPa

Air 896
H2, 0.1 MPa 786
H2, 1.03 MPa 703
H2, 6.89 MPa 662

* Tensile C, Appendix C-4.

- 89 -
DATA SHEET IA-16

Stress Necessary for Slow Crack Growth in

Type 304L Stainless Steel*

Net Section Time , hrs Crack

Stress, MN/m^ Incremental Accumu lated Growth

600 325 325 No

641 72 397 No
682 72 469 No
724 72 541 No
765 72 613 No
786 1.4 (failed) 614.4 Yes

* Crack developed during room temperature tensile test

in hydrogen environment; net section stress when
tensile test was stopped was 772 MN/m . Specimen
then loaded in creep frame at indicated stresses
without removal from the hydrogen environment.
Tensile E, Appendix C-5.

- 90 -
DATA SHEET IB-1
Type 304N Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yieldt Tensile Uniform Total Strain

298 Air none 760 880 33 1.24

Air 69 MPa H2** 740 830 31 1.05
69 MPa H2 none 640 840 36 0.78
69 MPa H2 69 MPa H2** 550 790 37 0.62
69 MPa He none 630 850 43 1.35

375 Air none 820 950 11 26 1.31

69 MPa D2tt 820 970 11 22 1.20

298 Air none 906 1110 16 28 1.47

69 MPa D2tt 950 1185 16 28 0.95

245 Air none 975 1340 27 37 1.82

69 MPa D2tt 1063 1420 22 27 0.49

220 Air none 1026 1450 26 35 1.67

69 MPa D2tt 1093 1480 21 24 0.33

200 Air none 1096 1810 47 56 1.44

69 MPa D2tt 1160 1510 19 23 0.38

* Tensile A, Appendix C-1; Heat Analysis, Appendix D-10.

** 69 MPa H2 at 430 K for 1000 hours.
t 0.2% offset,
tt 69 MPa D2 at 620 K for 3 weeks.

- 91 -
DATA SHEET IC-1

Type 309S Stainless Steel*

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Toti Strain

298 Air none 290 600 54 1.27

69 MPa He none 276 580 60 1.24
69 MPa H2 none 260 586 63 1.35
Air 69 MPa H2--430K 255 615 43 0.92
14d
Air 28 MPa H2--470K 330 615 57 1.17
100 hr

* Tensile A, Appendix C-1.

** 0.2% offset.

DATA SHEET IC-2

Tensile Properties of Type 309S Stainless Steels Containing

Hydrogen and Helium

Test (Condition Hydrogen Streng th, MPa Elongation,

Temp, K Environment Exposure Yield Ultimate %

300 Air none 243 612 56

300* Air ** 301 618 48-57
300* Air t 382 658 45-53
973tt Air none 131 296 27
973tt Air t 227 196 <1

* Specimens contained tritium and Helium-3.

** 328 mol hydrogen isotopes and 6.2 mol helium per m^ metal,
t 146 mol hydrogen isotopes and 2.5 mol helium per m metal,
tt Held 1/2 hour at 973 K before testing.

- 92 -
DATA SHEET ID-1

Type 310 Stainless Steel Bar Stock, As Received*

Test Condition Hydrogen Strength, MPa Elongation, % Fractu

Temp, K Environment Exposure** Yield Ultimate Uniform Total Strain

380 Air none 440 670 25 36 1.35

69 MPa 440 700 27 40 1.71

273 Air none 510 860 44 53 1.71

69 MPa 510 900 46 53 1.47
200 Air none 560 1200 60 66 1.20
69 MPa 590 1280 62 73 1.24
78 LN none 570 1720 74 78 1.05
69 MPa 570 1790 71 76 1.35

* Tensile B, Appendix C-2.

** Exposed at 470 K for 1449 days,

DATA SHEET ID-2

Type 310 Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fract

Temp, K Environment Exposure Yield** Tensile Uniform Total Straii

298 Air none 210 540 - 61 1.56

Air 69 MPa H2t 200 500 - 63 1.42
69 MPa H2 none 186 ,490 - 67 1.72
69 MPa H2 69 MPa H2t 180 440 - 66 1.56
69 MPa He none 180 480 — 70 1.61

* Tensile A, Appendix C-1.

** 0.2% offset,
t 69 MPa H2 at 430 K for 1000 hours

- 93 -
DATA SHEET IE-1

Type 316 Stainless Steel; Bar Stock, As Received*

Test Condition Hydrogen Strength, MPa Elongat ion, % Fracture

Temp, K Environment Exposure** Yield Ultimate Uniform Total Strain

none 810 830 7 20 1.61

380 Air 69 MPa 880 930 11 22 1.19

none 890 1040 21 33 1.47

273 Air 69 MPa 990 1160 20 32 1.13

none 900 1150 11 40 1.51

250 Air 69 MPa 1030 1280 Ik 35 1.07

none 960 1210 24 43 1.56

200 Air 69 MPa 1100 1410 26 37 1.06

* Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks.

- 94 -
DATA SHEET IF-1

Carpenter 20 Cb-3* Stainless Steel As-Received*

Test Condition Hydrogen Streng!th, MPa Elongation, % Fracture

Temp, K Environment Exposure** Yieldt Ultimate Uniform Total Strain

298 Air none 236 600 - 48 1.14

69 MPa H2 none 230 590 - 48 1.14
69 MPa H2 69 MPa D2 262 610 - 48 1.14

200 Air none 320 1100 60 66 1.01

Air 69 MPa D2 348 1177 55 62 1.08

* Tensile B, Appendix C-2; heat analysis. Appendix D-11.

** Exposed at 620 K for 3 weeks,
t 0.2% offset.

DATA SHEET IG-1

Incoloy* Alloy BOOH, Hot Rolled Plate, Solution Annealed*

Test Condition Hydrogen Strength, MPa Elongat ion, % Fracture

Temp, K Environment Exposure** Yield Ultimate Uniform Total Strain

380 Air none 270 750 47 52 0.92

69 MPa 290 770 48 53 0.82
298 Air none 310 820 48 55 0.97
69 MPa 330 840 48 53 0.78
250 Air none 340 870 48 56 1.20
69 MPa 360 900 49 55 0.92
200 Air none 360 930 49 56 1.11
69 MPa 380 990 54 63 0.94
78 LN none 530 1520 80 84 0.78
69 MPa 540 1490 74 76 0.69

* Tensile B, Appendix C-2; heat analysis. Appendix D-3.

** Exposed at 620 K for 3 weeks.

- 95 -
DAiyrAsmxTtSiiiri
T9^kUL6 28eeH.nle88 S t e e l i Bar S t o c k | 4 i Reeffived*

TflSBtetCcflixitttiTtfiiPn H^^idi2)iD[6n S8.liK5aiBgti\, MTOs

TfflqnpJCK E«nMMriioBmt<nt EiEpayngprt* YlfedBd^IIIttuiwtfea"
^x'**''' "^ ' -v
38098 AMr** n<mme 8810 ' SKSO -7 5^0 2\m
t o MPa
69 «T> He**
u .** 69none
MPa 1%° k%W _11 5'
273 Afi9 MPa H2** namme isao hWdO -21 513 0W«7
69 MPa 990 1160 20 32 1,13
298 Airt none 135 480 50 2.21
250 27
^ ^ MPa Het
69 MPa H2t
none
6^m •S530o
156 460
« ~24 4? 0.69
none 45
200 Air none 960 1210 24 43 1.56
69 MPa 1100 1410 26 37 1.06
* T e n s i l e A, Appendix C-1.
'**rdMsi£Mea, W^MM eti2tutes and furnace cooled.
** lE:^9«id*tt^ ^ilJSh^afidnealeide^d. K for 64 hours and a i r cooled.

DATA SHEET IH-2

Nickel 200, Notch-Bar T e n s i l e P r o p e r t i e s *

Test Condition Hydrogen St;rength, MPa Elongation, % Fracture

Temp, K Environment Exposure lield Ultimate Uniform Totiil Strain

298 Air** none - 660 - 0.35

69MPa He** none - 810 - 0.37
69 MPa H2** none - 560' - 0.11
Airt none - 635 - 0.44
69 MPa Het none - 710 - 0.34
69 MPa H2t none — 580 — — 0.20

* T e n s i l e A, Appendix C-1 w i t h notch.

** Annealed 1090 K 15 minutes and furnace cooled,
t As in **, plus annealed 773 K for 64 hours and a i r cooled.

- 96 -
DATA SHEET IJ-1

Nickel 301*

Test Condition Hydrogen Streng th, MPa Elongation, % Fracture

Temp,, K Environment Exposure Yield Ultimate Uniform Totial Strain

298 Air** none 451 778 - 39 1.89

69 MPa He** none 486 791 - 34 1.35
69 MPa H2** none 532 618 - 12 0.22

298 Airt none 1008 1380 - 23 0.49

69 MPa Het none 1009 1350 - 22 0.42
69 MPa H2t none — 850 — 4 0

* Tensile A, Appendix C-1.

** Annealed 1170 K for 5 min and quenched.
t Annealed as in **, plus annealed 860 K for 16 hours,
810 K for 5 hours and 755 K for 5 hours and furnace cooled.

DATA SHEET IJ-2

Nickel 301, Notch Bar T e n s i l e P r o p e r t i e s *

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield Ultimate Uniform Toti Strain

298 Air** none - 985 - 0.30

69 MPa He** none - 995 - 0.30
69 MPa H2** none - 690 - 0.01
Airt none - 1630 - 0.19
69 MPa Het none - 1600 - 0.10
69 MPa H2t none — 840 mm —
0.04

* T e n s i l e A, Appendix C-1 p l u s notch.

** Annealed 1170 K for 5 min and quenched.
t Annealed as in **, plus annealed 860 K for 16 h o u r s ,
810 K for 5 hours and 755 K for 5 hours and furnace cooled.

- 97 -
DATA SHEET IK-1

Type 440C Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure** Yield Ultimate Uniform Total Strain

298 Air none 377 620 7.1 0.010

69 MPa D2 377 575 4.6 0.006

200 Air none 406 670 7.7 0.013

69 MPa D2 450 570 4.2 0.009

* Tensile B, Appendix C-2.

** Exposure at 620 K for 3 weeks.

98 -
IRON-CHROMIUM-NICKEL-MANGANESE ALLOYS

DATA SHEET IIA-1

Tenelon* Plate, As Received*

Test Condition Hydrogen Streng th, MPa

Temp, K Environment Exposure** Yield Ultimate

350 Air none 675 1270

69 MPa 700 1300

273 Air none 830 1480

69 MPa 920 1540

200 Air none 1050 1960

69 MPa 1020 1620

78 LN none 1740 1780

nonet 1730 2040
nonett 1670 2120
nonel 1450 1730
69 MPa 1720 1780

* Tensile B, Appendix C-2; heat analysis. Appendix D

** Exposed at 620 K for 3 weeks.
t Electropolished.
tt Annealed 1170 K for 24 hours.
1 Annealed 1270 K for 24 hours.

- 99 -
DATA SHEET IIA-2

Tenelon**

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Tot< Strain

298 Air none 570 930 - 56 1.05

69 MPa He none 500 875 - 65 1.14
69 MPa H2 none 500 900 - 55 0.63
Air 69 MPa H2t 550 840 - 41 0.45
69 MPa H2 69 MPa H2t 470 760 — 24 0.26

* Tensile A, Appendix C-1.

** 0.2% offset,
t 69 MPa H2 for 1000 hours at 423 K.

DATA SHEET IIA-3

Fracture Toughness of Tenelon**

Test Specimen Fracture

Temp, K Condition Toughness, MPa Vm"

78 As received 68.6
Anneal 1170 K 36.5
Anneal 1270 K 71.4

200 As received 127.8

Anneal 1170 K 99.6
Anneal 1270 K 120.5

* Heat analysis. Appendix D-4; single edge

notched, Appendix C-6.

- 100 -
DATA SHEET IIB-1

Nitronic* 40 Stainless Steel Bar Stock, As Received*

Test Condition Hydrogen Strength, MPa Elongat ion, % Fracture

Temp, K Environment Exposure** Yield Ultimate Uniform Total Strain

none 680 940 30 39 1.66

380 Air 69 MPa 690 1020 36 46 1.02

none 770 1170 41 51 1.61

298 Air 69 MPa 800 1270 46 56 0.92

none 860 1360 46 57 1.51

250 Air 69 MPa - 1380 41 46 0.45

200 Air none 970 1550 48 58 1.56

69 MPa 1060 1650 44 48 0.65

78 LN none 1580 2140 45 49 0.64

69 MPa 1600 2060 36 36 0.38

* Tensile B, Appendix C-2; heat analysis. Appendix D-5.

** Exposed at 620 K for 3 weeks.

- 101 -
DATA SHEET IIB-2

Mechanical Properties of Nitronic* 40 Alloy: Heat Treatment and Notch Effects*

Test Strength, MPa Elongation, % Fracture

Temp, K** Treatment Specimen Yield Ultimate Uniform Total Strain

300 Solution Smooth 700 1170 41 51 1.59

Anneal bar
Notch 800 1160 24 27 0.74
bar
Sensitize Notch 750 1070 18 18 0.53
bar

200 Solution Smooth 880 1550 48 58 1.57

bar
Notch 1130 1500 19 19 0.72
bar
Sensitize Smooth 720 1490 51 60 1.03
bar
Notch 1120 1250 10 10 0.17
bar

* Heat analysis. Appendix D-5; Tensile B, Appendix C-2.

** Air environment.

- 102 -
DATA SHEET IIB-3

Mechanical Properties of Sensitized Nitronic* 40 Stainless Steel Tested

in a High-Pressure Hydrogen Environment at Room Temperature*

Specimen Test Strength, MPa Elongation, % Fracture

Condition Atmosphere Yield Ultimate Uniform Tot al Strain

Solution 69 MPa He 650 1050 42 52 1.11

Annealed
69 MPa H2 670 1060 41 50 1.22

920 K-2 hr 69 MPa He 640 1100 43 50 1.51

69 MPa H2 640 1080 42 49 1.50
920 K-24 hr 69 MPa He 625 1110 46 53 1.38
69 MPa H2 620 1100 46 52 1.10
920 K-24 hr 69 MPa He** 760 1060 16 16 0.32
60 MPa H2** 700 760 9 9 0.06

* Heat analysis, Appendix D-5; Tensile B, Appendix C-2.

** Notch bar specimen.

DATA SHEET IIB-4

Mechanical Properties of Sensitized Nitronic* 40 Stainless Steel

Saturated with Hydrogen*

Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Treatment** Exposure Yield Ultimate Uniform Tot;al Strain

200 Solution none 970 1550 48 58 1.57

Anneal
Solution 69 MPa H2t 1060 1650 44 48 0.66
Anneal
920 K-24 hr none 790 1490 51 60 1.03
920 K-24 hr 69 MPa H2 920 1470 37 37 0.33
920 K-24 hr 69 MPa H2tt 900 1350 35 40 0.45

* Heat analysis. Appendix D-5; Tensile B, Appendix C-2.

** Smooth bar tensile specimens.
t Exposed at 620 K for 3 weeks,
tt Crosshead speed, 5 mm/sec; all others, 0.5 mm/sec.

- 103 -
DATA SHEET IIB-5

Nitronic • 40*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 400 670 58 1.51

69 MPa He none 350 700 59 1.47
69 MPa H2 none 360 700 61 1.43

* Tensile A, Appendix C-1.

** 0.2% offset.

DATA SHEET IIB-6

Nitronic* 40; Cold Worked 30Z*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 1240 1290 _ 26 0.87

30 MPa H2 1075 1150 - 32 0.43
69 MPa He none 1010 1050 - 26 0.99
69 MPa H2 none 980 1100 - 26 1.02
69 MPa H2 30 MPa H2 1060 1130 — 36 0.44

* Tensile A, Appendix C-1.

** 0.2% offset.

- 104 -
DATA SHEET IIB-7

Nitronic* 40 Stainless Steel, High Energy Rate Forged*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield Ultimate Uniform Total Strain

380 Air none 540 1040 47 59 1.81

69 MPa** 570 1070 50 68 1.26
none 780 970 21 31 1.17
69 MPat 690 930 26 33 0.67
298 Air none 780 1140 32 44 1.24
69 MPat 890 1220 30 42 0.96
273 Air none 640 1300 57 69 1.81
69 MPa** 690 1430 67 78 1.06
220 Air none 900 1320 33 45 1.31
69 MPat 960 1420 37 47 0.80
200 Air none 930 1700 51 59 1.26
69 MPa** 1050 1830 49 54 0.90
none 1020 1610 42 54 1.26
69 MPat 990 1740 53 60 0.66
78 LN none 1450 2840 46 56 0.83
69 MPa** 1400 2600 46 46 0.53

* Tensile B, Appendix C-2.

** 69 MPa at 470 K for 1449 days,
t 69 MPa at 620 K for 21 days.

DATA SHEET IIB-8

Nitronic* 40 Stainless Steel, High Energy Rate Forged*

Test Condition Hydrogen Streingth , MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uni form Tot;!i Strain

298 Air none 610 790 - 34 1.35

28 MPa H2 660 820 - 31 0.89
69 MPa He none 570 780 - 34 1.39
69 MPa H2 none 570 790 - 30 1.31
28 MPa H2 630 830 — 31 0.78

* Tensile A, Appendix C-1.

** 0.2% offset.

- 105 -
DATA SHEET IIB-9

Nitronic* 40 Stainless Steel, High Energy Rate Forged*

Test Condition Hydrogen Impact

Temp, K Environment Exposure Energy, J

298 Air none 110

29.6 MPa H2** 91

77 LN none 37
29.6 MPa H2** 35

* Impact, Appendix C-8.

** 470 K - 56 days.

DATA SHEET IIB-10

Nitronic* 40 Stainless Steel, High Energy Rate Forged*

Fracture Toughness,
Test Condition Hydrogen MPa "Vm

Temp, K Environment Exposure Long Trans

298 69 MPa He none 79 74

69 MPa H2 none 81 68

69 MPa H2 0.6 MPa H2 76 62

* C-shaped tensile. Appendix C-7.

- 106 -
DATA SHEET IIB-11

Fracture Parameters for N i t r o n i c * 40 S t a i n l e s s S t e e l *

1000 —
I I Tested in Helium

y/j^ Tested in Hydrogen

[ I Charged with Deuterium, Tested in Hydrogen

800
r^^y Annealed in Helium, Tested in Hydrogen

t 600

T3
01
•(->
m

01 C
400
1— 01
^ oo
O =^ .

200
^'.
rt
ft
/ k
Parallel 45° 90°
Orientation Orientation Orientation

* C-shaped t e n s i l e , Appendix C-7. Test i n 69 MPa He or H2.

Deuterium exposed at 69 MPa at 620 K for 3 weeks.

- 107 -
DATA SHEET IIB-12
Fracture Parameters for Nitronic* 40 Stainless Steel*

1 1 Tested i n Helium 1 ] Charged with Deuterium,

1 1 1—,—i Tested i n Hydrogen
l/A Tested in Hydrogen fvH'l Annealed in Helium, Tested 1
500 —L_ZJ ^ ^ L j i J in Hydrogen

T
•
-

400
T —
— -
* •• _
7 7
/ /
n3
300 —
/ /
//
/ /
A
i 7 "TT-

1
/
T3 /
T3 /
/
V,

i
V
200

••'.*'••
1
/ y •••;;"•

100 — // -H

i
n %. ' ^ ^ / •• • • .

Parallel 45° 90°

Orientation Orientation Orientation

* C-shaped t e n s i l e . Appendix C-7. Test in 69 MPa He or H2.

Deuterium charged a t 69 MPa a t 620 K for 3 weeks.

- 108 -
DATA SHEET IIC-1

Nitronic* 50 Stainless Steel Bar Stock, As Received*

Test Condition Hydrogen Strength, MPa Elongat ion, % Fracture

Temp J
, K Environment Exposure** Yield Ultimate Uniform Total Strain

380 Air none 700 990 26 34 1.23

69 MPa 720 1060 29 37 1.16
298 Air none 800 1190 32 41 1.17
69 MPa 820 1240 33 44 1.05
248 Air none 870 1310 34 43 1.23
69 MPa 900 1390 35 43 1.00
200 Air none 1030 1550 35 44 1.08
69 MPa 1020 1620 37 44 0.97
78 LN none 1590 2310 38 44 0.91
69 MPa 1590 2350 38 44 0.90

* Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks,

DATA SHEET IIC-2

Nitronic* 50 Stainless Steel Bar Stock, As Received.*

Test Condition Hydrogen Strength, MPa Fracture

Temp, K Environment Exposure Yield** Tensile Elomgat ion, 3: Strain

298 Air none 440 710 43 1.27

69 MPa He none 400 680 47 1.35

•
69 MPa H2 none 400 680 45 1.31

* Tensile A, Appendix C-1; heat analysis. Appendix D-6.

** 0.2% offset.

- 109 -
DATA SHEET IIC-3

Nitronic* 50 Stainless Steel, High-Energy-Rate-Forged*

Test Condition Hydrogen Deflection J^ Jm dJ/da

Temp,, K Environment Exposure mm kJ/m2 MPa

298 69 MPa Het none - 32 176

69 MPa H2 no^ne - 23 137
69 MPa H2 D2 - 33 211

69 MPa Hett none - 936 360

69 MPa H2tt none - 107 209
69 MPa H2tt D2 181 264

* C-Shaped tensile. Appendix C-7.

** Exposed at 620 K for 3 weeks.
t Crack parallel to forging pattern
tt Crack perpendicular to forging patterns,

DATA SHEET IID-1

18-18 Plus* Stainless Steel*

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Exposure Yield** Tensile Uniform Total Strain

298 69 MPa He none 520 910 63 1.51

69 MPa H2 none 506 880 42 0.42

* Tensile A, Appendix C-1; heat analysis. Appendix D-9.

** 0.2% offset.

- 110 -
DATA SHEET IIE-1

X18-3 Mn Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 580 810 - 45 1.24

69 MPa He none 530 790 - 50 1.35

69 MPa H2 none 520 790 - 46 1.31

* Tensile A, Appendix C-1; heat analysis. Appendix D-8.
** 0.2% offset.

DATA SHEET IIF-1

18-2 Mn Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain
298 Air none 730 1007 - 51 0.87

69 MPa H2 none 660 924 - 33 0.31

* Tensile A, Appendix C-1.

** 0.2% offset.

- HI -
DATA SHEET IIG-1

Type 216 Stainless Steel*

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Temp J, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 640 810 40 1.10

298 Air 69 MPa H2t 630 790 36 1.05

298 69 MPa H2 none 590 780 44 1.17

298 69 MPa H2 69 MPa H2t 560 760 45 1.02

298 69 MPa He none 590 790 45 1.20

* Tensile A, Appendix C-1; heat analysis. Appendix D-7.

** 0.2% offset,
t 69 MPa H2 at 430 K for 1000 hours.

- 112 -
PRECIPITATION HARDENABIZ ALLOYS

DATA SHEET IIIA-1

A-286 Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield Tensile Uniform Total Strain

298 Air none 765 1098 - 25 0.77

2.1 MPa Argon 776 1089 - 24 0.79
69 MPa D/T** 750 1041 - 13 0.34
nonel - 1500 - 4 0.15
2.1 MPa Argoni - 1380 - 3.6 0.11
69 MPa D/T**.1 - 1310 - 3 0.06
none lOlOt 1350tt 23 28 0.50
69 MPa D2II 1070t 1380tt 23 24 0.24

220 Air none llOOt 1520tt 28 34 0.49

69 MPa D2II 1130t 1530tt 27 27 0.25

* Tensile A, Appendix C-1.

** 69 MPa D/T at 370 K for 200 days.
t True stress at 5% strain,
tt True stress at maximum load.
1 Notched-bar tensile specimens, all others smooth-bar specimens.
11 69 MPa D2 at 620 K for 3 weeks.

- 113 -
DATA SHEET IIIA-2

A-286 St ainless Steel High Energy, Rate Forged*

Fracture
Test Condition Hydrogen Toughness,
Temp, K Environment Exposure MPa Vl"

298 69 MPa He none 76**

69 MPa H2 89**
69 MPa He none 71***
69 MPa H2 90***
69 MPa He none 81t
69 MPa H2 82t
69 MPa He none 93tt
69 MPa H2 89tt
69 MPa He 1.6 MPa D2 88tt
69 MPa H2 1.6 MPa D2 97tt
69 MPa He none 521
69 MPa H2 none 561
69 MPa H2 1.5 MPa D2 591
69 MPa He none 9311
69 MPa H2 none 9011
69 MPa H2 1.5 MPa D2 9711

* Single edge notched. Appendix C-6.

** Aged 4 hours at 990 K (Heat 1).
*** Aged 8 hours at 990 K (Heat 1).
t Aged 16 hours at 990 K (Heat 1).
tt Aged 8 hours at 990 K (Heat 2).
1 HERF only not aged. Rc-H-
11 Aged 8 hours at 990 K. Rc-H-

- 114 -
DATA SHEET IIIA-3

A-286 Stainless Steel Notch Impact Test*

Test Condition Hydrogen Impact

Temp, K Environment Exposure Energy, J

298 Air Base Metal

As Received 6,10
Argon** 5,08
D/Tt 4.74

298 Air Weld Metal

As Received 4.18
Argon** 3.40
D/Tt 4.51

* Impact Appendix C-8.

** 0.21 MPa at 370 K for 200 days,
t 69 MPa D/T at 370 K for 200 days.

- 115 -
DATA SHEET IIIA-4

Fracture Parameters for A-286 Stainless Steel*

1000 —
Tested in Helium
1 ' ' II
/ / Tested in Hydrogen

Charged with Deuterium, Tested in Hydrogen

800 —

600 ^^^
—

E
-
q
/
(,
400 / y

Y/ / nr
200 — :•
-r

-±-
1
Parallel
Orientation
*1
-55°
'A
Orientation
A '

90°
Orientation

* C-shaped t e n s i l e . Appendix C-7. T e s t e d in 69 MPa H2 or He.

Deuterium charged a t 69 MPa a t 620 K for 3 weeks.

- 116 -
DATA SHEET IIIA-5

Fracture Parameters for A-286 Stainless Steel*

Tested in Helium

. '/^ Tested in Hydrogen

400 — Charged with Deuterium, Tested in Hydrogen —

.,r
300
~7 1r T
-?
Q.
• •

^ r — R —
yY
T3
-L
y<
• / A
J
Y/ '
A •
y
Y/,
^ /
/Y' Y Y Y

200 Y/. YA / /

V/
V/
VY
YY
y
Y/y
100
A.\J\J
T Y' X
X X

1
^
'

y//
V/
'// YY
n YA \// v/
Parallel 45° 90°
Orientation Orientation Orientation

* C-shaped T e n s i l e , Appendix C-7. Tested in 69 MPa He or H2.

Deuterium charged a t 69 MPa a t 620 K for 3 weeks.

- 117 -
DATA SHEET IIIB-1

JBK-75 HERF and Age*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 69 MPa He none 800 1090 10 14 0.63

69 MPa H2 none 890 1160 10 13 0.40

* Tensile, Appendix C-3.

** 0.2% offset.

DATA SHEET IIIB-2

JBK-75 HERF and Age*

Stress Fracture
Test Condition Hydrogen Intens ity. Energy,
Temp, K Environment Exposure MPa V¥ MJ/m^

298 69 MPa He none 80 0.350

69 MPa H2 none 80 0.333
69 MPa H2 0.7 MPa D2 81 0.294
at 625 K

* c-shaped tensile. Appendix C-7.

- 118 -
DATA SHEET IIIC-1

17-4 Stainless Steel, Tensile Tubes*

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 940 980 4.7 -

69 MPa He 69 MPa He 1076 1145 6.4 -
35 MPa D/T 35 MPa D/Tt 1000 1000 0.7 -
69 MPa D/T 69 MPa D/Ttt 1062 1096 1.2 —

* Tensile E, Appendix C-5.

** 0.2% offset.
t 8 hours at 315 K.
tt 2 hours at 370 K.

DATA SHEET IIIC-2

Fracture Toughness 17-4 PH Stainless Steel*

Fracture Toughness, M P a V ^

Material Test Environment

Condition 69 MPa He 3.5 MPa H? 69 MPa H,

Underaged 104 V 31 20
Peak Aged 97 29 13
Overaged - 57 34
Solution Annealied 97 71 31

Heat Treatments

Material Aging Hardness

Condition Temp, K %i
Underaged 709 38
Peak Aged 783 42
Overaged 866 35
Solution Annealed ~ 28

* c-shaped tensile. Appendix C-7.

All specimens were solution annealed 2 hours at
1339 K and aged 1 hour at indicated temperatures,

- 119 -
DATA SHEET IIID-1

AM-350 Stainless Steel*

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 420 1160 - 70 -

69 MPat 455 580 - 3/4 -
69 MPa He none 420 1240 - 55 -
6.9 MPa D2 none 345 430 - 4 -
69 MPa D2 none 430 520 - 2.6 -
0.69 MPa D2 none 410 455 ~ 3 ~

* Condition H - annealed at 1310 to 1350 K air cool or water quench,

** 0.2% offset,
t 26 days at 570 K.

DATA SHEET IIIE-1

AM-363 Stainless Steel

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield* Tensile Uniform Total Strain

298 Air none 890 890 7 -

Air 0.21 MPa H2** 900 900 8.6 -
Air none 134at 1480 3 -
Air 0.21 MPa H2** 1400t 1500 3 • "

* 0.21 MPa D2 at 630 K for 5 days.

** 0.2% offset,
t Notched - 45° notch. Notch diameter = 0.5X outer diameter.

- 120 -
DATA SHEET IIIF-1

C6-27 Stainless Steel*

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 69 MPa He none 806 1165 29 0.30

69 MPa H2 none 855 1117 10 0.13
298 69 MPa H2 69 MPa H2 at 855 1020 4 0.03
425 K-72 hrs
298 69 MPa Het none 1070 1385 12 0.13
69 MPa H2t none 1034 1138 1 0.03

* Tensile A, Appendix C-1

** 0.2% offset,
t HERF specimens.

DATA SHEET IIIG-1

Ni-SPAN-C* (Alloy 902)

Test Condition Hydrogen Strength , MPa Elongation, % Fracture

Temp, K Environment Exposure Yield** Tensile Uniform Total Strain

298 Air none 676 1186 - 10 -

69 MPa He none 750 1160 - 16 -
6.9 MPa H2 none - 1170 14
69 MPa H2 none 650 1130 - 15 -

* Sheet specimens 0.25 mm and 19 mm gauge length.

** 0.2% offset.

- 121 -
HIGH PURITY ALLOYS

DATA SHEET IVA-1

Mechanical Properties (Alloy A ) *

Test Condition Hydrogen Strength, MPa Elongation, % Fracture

Temp, K Environment Exposure** Yield Ultimate Uniform Total Strain

370 Air none 230 610 45 52 1.57

69 MPa 270 660 50 59 1.65
298 Air none 350 1270 62 73 1.66
69 MPa 290 1030 60 60 0.50
235 Air 69 MPa 390 1110 38 38 0.27
200 Air none 540 1320 36 46 1.42
69 MPa 420 1190 33 33 0.31
78 LN none - - - - 1.44
69 MPa 1060 42 48 1.13

* Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks.

- 122 -
DATA SHEET IVB-1

Mechanical Properties (Alloy B ) *

Test Condition Hydrogen Strength, MPa Elongat ion, % Fractu

Temp,, K Environment Exposure** Yield Ultimate Uniform Total Strain

370 Air none 240 630 45 56 1.58

69 MPa 260 660 46 56 1.40
298 Air none 340 1020 61 69 1.56
69 MPa 290 870 65 72 1.50
235 Air 69 MPa 320 1170 72 79 0.44
200 Air none 340 1170 64 74 1.57
69 MPa 380 1250 66 71 0.89
78 LN none 260 870 63 67 1.37
69 MPa 270 900 66 72 1.41

* Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks.

DATA SHEET IVC-1

Mechanical Properties (Alloy C ) *

Test Condition Hydrogen Strength, MPa Elongation. % Fractu

Temp J
, K Environment Exposure** Yield Ultimate Uni form Totial Strain

370 Air none 250 630 44 52 1.62

69 MPa 260 660 45 53 1.45
298 Air none 330 910 49 58 1.65
69 MPa 290 770 52 62 1.55
200 Air none 300 1100 78 87 1.52
69 MPa 330 1170 78 86 1.50
78 LN none 250 850 82 89 1.53
69 MPa 280 890 80 86 1.43

* Tensile B, Appendix C-2.

** Exposed at 620 K for 3 weeks.

- 123 -
REFERENCES

1. I. M. Bernstein and A. W. Thompson. Hydrogen in Metals, ed.,

American Society for Metals, Metals Park, Ohio (1974).

2. A. W. Thompson and I. M. Bernstein. Effect of Hydrogen on

Behavior of Materials, ed.. Metallurgical Society of AIME, New
York, NY (1976).

3. A. W. Thompson and I. M. Bernstein. Effect of Hydrogen on

Behavior of Materials, ed.. Metallurgical Society of AIME,
Warrendale, PA (1981).

4. M. R. Louthan, Jr. and R. P. McNitt. Environmental Degradation

of Engineering Materials, ed., Virginia Polytechnic Institute,
Blacksburg, VA (1977).

5. M. R. Louthan, Jr., R. P. McNitt, and R. D. Sisson, Jr.

Environmental Degradation of Engineering Materials - II, ed.,
Virginia Polytechnic Institute, Blacksburg, VA (1981).

6. MM J. P. Fidelle et. M. Papin. L' Hydrogen Dans Les Metaux,

Colloque - Valduc (27-28, September 1967).

7. "L'Hydrogen Dans Les Metaux-2." Congress International, Paris

(29 Mai-2 Juin, 1972).

8. "Hydrogen in Metals." Proceedings of the 2nd International

Congress, Paris, 6-10 June 1977. Pergamon Press, New York,
(1977).

9. "Hydrogen in Metals," Proceedings of the 3rd International

Congress, Paris, 7-2 June 1982.

10. "Hydrogen Embrittlement Testing." ASTM Special Technical

Publication 543, American Society for Testing and Materials,
Philadelphia, PA (1974).

11. Advanced Techniques for Characterizing Hydrogen in Metals,

TMS-AIME, Warrendale, PA (1982).

12. C D . Beachem, ed. Hydrogen Damage, American Society for

Metals, Metals Park, OH (1979).

- 124 -
J. p. Hirth and H. H. Johnson. "Hydrogen Problems in Energy
Related Technology." Corrosion 32, 3 (1976).

J. J. Au and H. K. Birnbaum. "Hydrogen Effects on the

Deformation of Iron Whiskers." Scripta Met 12, 457 (1978).

G. C. Smith. "Effect of Hydrogen on Nickel and Nickel-Base

Alloys." Hydrogen in Metals, p. 485, American Society for
Metals, Metals Park, Ohio (1974).

J. K. Tien, A. W. Thompson, I. M. Bernstein, and Rebecca J.

Richards. "Hydrogen Transport by Dislocations." Met. Trans.
7A, 821 (1976).

A. W. Thompson, "The Mechanism of Hydrogen Participation in

Ductile Fracture." Effect of Hydrogen on Behavior of
Materials, p. 467, Metallurgical Society of AIME, Warrendale,
PA (1976).

J. F. Breedis and L. Kaufmann. "Formation of HCP and BCC

Phases in Austenitic Iron Alloys." Met. Trans. 2, 2359 (1971).

D. V. Neff, T. E. Mitchell, and A. R. Troiano. "The Influence

of Temperature, Transformation, and Strain-Rate on the
Ductility Properties of Austenitic Stainless Steels." Trans.
A ^ 62, 858 (1969).

B. C. Odegard, J. A. Brooks, and A. J. West. "The Effect of

Hydrogen on the Mechanical Behavior of Nitrogen Strengthened
Stainless Steel." p. 116 in Effect of Hydrogen on Behavior of
Materials. Metallurical Society of AIME, Warrendale, PA (1976),

R. E. Stoltz and J. B. Vander Sande. "The Effect of Nitrogen

on Stacking Fault Energy of Fe-Cr-Ni-Mn Steels." Met. Trans.
IIA, 1033 (1980).

H. Mecking. "Deformation of Polycrystals." p. 1573 in

Strength of Metals and Alloys, Vol. 3, Pergamon Press, New
York (1980).

C. L. Briant. "Hydrogen Assisted Cracking of Sensitized 304

Stainless Steel." Met. Sci. 9A, 731 (1978).

A. W. Thompson. "The Behavior of Sensitized 309S Stainless

Steel in Hydrogen." Mat. Sci. and Eng. 14, 253 (1974).

J. D. Hobson and J. Hewitt. "The Effect of Hydrogen on the

Tensile Properties of Steel." Journal Iron and Steel
Institute 172, 131 (1953).

- 125 -
H. R. Gray. "Testing for Hydrogen Environment Embrittlement:
Experimental Variables." p. 133 in Hydrogen Embrittlement
Testing. ASTM STP 543, American Society for Testing and
Materials, Philadelphia, PA (1974).

M. B. Whiteman and A. R. Troiano. "Hydrogen Embrittlement of

Austenitic Stainless Steels." Corrosion 21, 53 (1965).

J. P. Fidelle, et al. "Disk Pressure Testing of Hydrogen

Environment Embrittlement." p. 221 in Hydrogen Embrittlement
Testing. ASTM STP 543, American Society for Testing and
Materials, Philadelphia, PA (1974).

D. Eliezer, D. G. Chakrapani, C. J. Alstetter, and E. N. Pugh.

"The Influence of Austenite Stability on the Hydrogen Embrit-
tlement and Stress Corrosion Cracking of Stainless Steel."
Met. Trans. lOA, 935 (1979).

D. Eliezer. Hydrogen Assisted Cracking in Type 304L and 316L

Stainless Steel. Third International Conference on Effect of
Hydrogen on Behavior of Materials, August 1980.

R. J. Walter and W. T. Chandler, p. 104 discussion to paper by

A. W. Thompson. "Ductility Losses in Stainless Steel."
Hydrogen in Metals. American Society for Metals, Metals Park,
Ohio (1974).

M. W. Perra. "Sustained-Load Cracking of Austenitic Steels in

Gaseous Hydrogen." Environmental Degradation of Engineering
Materials, ed., M. R. Louthan, Jr.,( R. P. McNitt, and R. D.
Sisson, Jr., VPI, Blacksburg, VA 321-333 (1981).

A. W. Loginow and E. H. Phelps "Steels for Seamless Hydrogen

Pressure Vessels" Corrosion 31 404 (1975).

R. L. Tobler and R. P. Reed. "Tensile and Fracture Behavior

of a Nitrogen Strengthened Chromium-Nickel-Manganese Stainless
Steel at Cryogenic Temperatures", p. 537 in Special Technical
Publication, ASTM STP 668. American Society for Testing and
Materials, Philadelphia, PA (1979).

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

Relationship for Interstitially Strengthened AISI 304 Stain-
less Steel at 4 K Temperature," p. 250 in Special Technical
Publication ASTM STP 743, American Society for Testing and
Materials, Philadelphia, PA (1982).

W. Losch and I. Andreoni. "Grain Boundary Surface Phases

Observed in Temper Embrittled High-Alloy Steels." Scripta
Met. 12, 277 (1978).

- 126 -
A. W. Thompson. "Mechanical Behavior of Face-Centered Cubic
Metals Containing Helium." Mater. Sci. and Eng. 21, 41-48
(1975).

M. R. Louthan, Jr., G. R. Caskey, D. E. Rawl, Jr., and C. W.

Krapp. "Tritium Effects in Austenitic Steels." p. 117 in
Proceedings of International Conference on Radiation Effects
and Tritium Technology for Fusion Reactors IV. CONF-752026-13
(1976).

D. E. Rawl, Jr., G. R. Caskey, Jr., and J. A. Donovan. "Low

Temperature Helium Embrittlement of Tritium Charged Stainless
Steel." 109th Annual AIME Meeting, Las Vegas, Nevada, February
24-28, 1980.

J. A. Donovan. "Effects of Helium on Mechanical Properties of

Armco Iron." Fall TMS-AIME Meeting, Niagara Falls, New York,
September 20-23, 1976.

A. J. West and D. E. Rawl, Jr. "Hydrogen in Stainless Steel:

Isotopic Effects on Mechanical Properties." Proceedings of
Conference in Tritium Technology in Fusion, Fusion and
Isotopic Applications. Conference 800427 (1980).

M. R. Louthan and R. G. Derrick, "Hydrogen Transport in

Austenitic Stainless Steel." Corros. Sci. 15 565 (1975).

G. R. Caskey, Jr., and R. D. Sisson, Jr. "Hydrogen Solubility

in Austenitic Stainless Steel." Scripta Met 15, 1187 (1981).

M. R. Dietrich, G. R. Caskey, Jr., and J. A. Donovan.

J-Controlled Crack Growth as an Indicator of Hydrogen-Stainless
Steel Compatibility, Hydrogen Effects in Metals, ed.
I. M. Bernstein and A. W. Thompson, Metallurgical Society of
AIME, Warrendale, PA, p. 637 (1981).

G. R. Caskey, Jr., D. A. Mezzanotte, Jr., and D. E. Rawl, Jr.

"Helium Embrittlement Stainless Steels at Ambient Temperature."
Scripta Met 16. 969 (1982).

T. L. Capeletti and M. R. Louthan, Jr. "The Tensile Ductility

of Austenitic Steels in Air and Hydrogen." Journal of
Engineering Materials and Technology, 99, 153-158 (1977).

T. L. Capeletti. "Effect of Hydrogen Exposure or the Fracture

Toughness of 17-4 pH Stainless Steel." Int. Conference on
Mechanical Behavior of Materials, Boston, August 16-20 (1976).
Conference 760819-12.

- 127 -
G. R. Caskey, Jr. "The Role of Twinning and Transformation in
Hydrogen Embrittlement of Austenitic Stainless Steels" in
Environmental Degradation of Engineering Materials, ed. by
M. R. Louthan and R. P. McNitt, Virginia Polytechnic Institute,
Blacksburg, VA, pp. 437-449 (1977).

G. R. Caskey, Jr. "Fractography of Hydrogen Embrittled

Stainless Steel." Scripta Met 11, 1077-1083 (1977).

G. R. Caskey, Jr. "A Mechanism for Hydrogen-Induced Brittle

Fracture of Austenitic Stainless Steel." DP-MS-77-86,
E. I. du Pont de Nemours & Co., Savannah River Laboratory,
Aiken, SC, 107th Annual AIME Meeting, Denver, CO, February 26 -
March 2, 1978.

G. R. Caskey, Jr. "Fracture of Fe-Cr-Mn Austenitic Steel."

DP-MS-78-68, E. I. du Pont de Nemours & Co., Savannah River
Laboratory, Aiken, SC, 108th Annual AIME Meeting, New Orleans,
LA, February 18-20, 1979.

G. R. Caskey, Jr. and J. A. Donovan. "Effect of Heat Treatment

on Fracture of Type 304L Stainless Steel in a Hydrogen Environ-
ment." DP-MS-79-18, E. I. du Pont de Nemours & Co., Savannah
River Laboratory, Aiken, SC, 1979, TMS-AIME Fall Meeting,
Milwaukee, WI, September 16-20, 1979.

G. R. Caskey, Jr. "Notch and Hydrogen Effects on Sensitized

21-6-9 Stainless Steel." DP-MS-79-22, E. I. du Pont de Nemours
& Co., Savannah River Laboratory, Aiken, SC, 1979 TMS-AIME Fall
Meeting, Milwaukee, WI, September 16-20, 1979.

G. R. Caskey, Jr. "Hydrogen Effects on Flow Localization and

Fracture in Type 304L Stainless Steel." DP-MS-79-57,
E. I. du Pont de Nemours & Co., Savannah River Laboratory,
Aiken, SC, Gordon Research Conference, Andover, NH, July 9-13,
1979.

G. R. Caskey, Jr. "Acoustic Emission from Hydrogen Saturated

Type 304L Stainless Steel." Scripta Met 13 583-587 (1979).

G. R. Caskey, Jr. "Hydrogen Assisted Fracture of Sensitized

Type 304L Austenitic Stainless Steel." DP-MS-79-62,
E. I. du Pont de Nemours & Co., Savannah River Laboratory,
Aiken, SC, 109th Annual AIME Meeting, Las Vegas, NV, February
24-28, 1980.

G. R. Caskey, Jr. "Fractography of Hydrogen-Embrittled Iron

Chromium-Nickel Alloys." Microstructural Science £, 413-419
(1981).

- 128 -
G. R. Caskey, Jr. "Hydrogen Induced Brittle Fracture of Type
304L Austenitic Stainless Steel." American Society for Testing
and Materials, STP 733, 86-97 (1981).

G. R. Caskey, Jr. "Effect of Hydrogen on Work Hardening of

Type 304L Austenitic Stainless Steel." Scripta Met 15,
1183-1186 (1981).

G. R. Caskey, Jr. "Hydrogen Damage in Stainless Steel."

Environmental Degradation of Engineering Materials. Edited by
M. R. Louthan, Jr., R. P. McNitt, and R. D. Sisson, Virginia
Polytechnic Institute, Blacksburg, VA, pp. 283-302 (1981).

G. R. Caskey, Jr. "Hydrogen Effects in High Purity Stainless

Steels." Third International Congress on Hydrogen and
Materials, Paris, France, June 7-11, 1982.

J. A. Donovan. "Accelerated Evolution of Hydrogen from Metals

During Plastic Deformation." Met. Trans. 7A, 1677-1683
(1977).

M. L. Holzworth and M. R. Louthan, Jr. "Hydrogen Induced Phase

Transformations on Type 304L Stainless Steels." Corrosion 24,
110-124 (1968).

M. L. Holzworth. "Hydrogen Embrittlement of Type 304L

Stainless Steel." Corrosion 25, 107-115 (1969).

M. R. Louthan, Jr., G. R. Caskey, Jr., J. A. Donovan, and

D. E. Rawl, Jr. "Hydrogen Embrittlement of Metals." Mater.
Sci. and Eng. 10 357 (1972).

M. R. Louthan, Jr., J. A. Donovan, and D. E. Rawl, Jr. "Effect

of High Dislocation Density on Stress Corrosion Cracking and
Hydrogen Embrittlement of Type 304L Stainless Steel."
Corrosion 29, 108-111 (1973).

M. R. Louthan, Jr., J. A. Donovan, and G. R. Caskey, Jr.

"Tritium Absorption in Type 304L Stainless Steel." Nuclear
Technology 26, 192-200 (1975).

M. R. Louthan, Jr., D. E. Rawl, Jr., J. A. Donovan, and

W. G. Holmes. "Hydrogen Embrittlement of Austenitic Stainless
Steels." Trans. Am. Nuc. Soc. 21, 158 (1975).

M. R. Louthan, Jr. and G. R. Caskey, Jr. "Hydrogen Transport

and Embrittlement in Structural Metals." 1st World Hydrogen
Energy Conference (1976).

- 129 -
M. R. Louthan, Jr. and R. P. McNitt. "The Role of Test
Technique in Evaluating Hydrogen Embrittlement Mechanism."
Effect of Hydrogen in Behavior of Materials, ed. A. W. Thompson
and I. M. Bernstein, Metallurgical Society of AIME, Warrendale,
PA, p. 496 (1976).

M. R. Louthan, Jr. "Effect of Hydrogen on the Mechanical

Properties of Low Carbon and Austenitic Steels." Hydrogen in
Metals, ed. T. M. Bernstein and A. W. Thompson, American
Society for Metals, Metals Park, Ohio, pp. 53-77 (1974).

- 130 -
APPENDIX A. DEFINITIONS

ELONGATION

Percentage increase of a gauge length, usually one inch,

during plastic strain in tension. In the data presented here,
crosshead motion was taken as the measure of change in length.

Total elongation is length increase at fracture.

Uniform elongation is length increase to the point where drop

in load is detected which signals beginning of observable strain
localization or necking.

HEAT TREATMENTS

Aging is a process of heating a previously solution-treated

alloy to an intermediate temperature to cause precipitation of a
finely dispersed phase which hardens the alloy.

Sensitization is a heat treatment that causes precipitation of

carbides of the form M23 C5 along grain boundaries and simulta-
neously reduces the chromium content of the grain boundary regions.

Solution annealing is a process of heating to elevated temper-

ature to dissolve all precipitates and produce a homogeneous solid
solution and quenching to retain the solid solution.

MECHANICAL PROCESSING

Ingots of stainless steel are formed into plate or bar by

mechanical processes of rolling and forging.

Cross-rolled plate refers to turning plate 90° between passes

through the rolling mills to minimize preferred orientation that
arises during the rolling process.

High energy rate forged ( H E R F ) alloys are hot forged at a very

rapid rate and immediately quenched in water to retain deformation
introduced during forging.

- 131 -
PLASTIC STRAIN

Irreversible or permanent strain of the test specimen measured

by subtracting elastic or recoverable strain from total strain.
This was usually done graphically on the load-deformation record
obtained during a tensile test.

Plastic strain to failure ( _) is calculated from the mea-

sured change in cross sectional area from the original (AQ) to
the final area (Af) at the fracture.

^p = In Ao/Af

Reduction in area (RA) is a measure of plasticity calculated from

the original (AQ) and final (Af) cross sectional areas.

RA = 100 — 4 —
^0

STRESS

Stress or force per unit area may be defined with respect to

an initial area (engineering stress) or the instantaneous area
(true stress). Both definitions have been utilized in data
presented here and are distinguished in each table.

Yield strength is the stress corresponding to a plastic strain

of 5% unless otherwise noted.

Ultimate strength is the true stress corresponding at maximum

load.

Tensile strength is the engineering stress at maximum load.

STRESS INTENSITY

The stress intensity factor (K) relates the stress field

(o£j) around a crack tip to the crack dimensions (a) and
specimen dimensions (width = w ) , where the function f(a,W) depends
on specimen shape, crack location and loading mode.

The stress intensity corresponding to the critical value for

crack extension is the Fracture Toughness (K^). Fracture tough-
ness is a measure of the ability of a material to resist crack
propagation.

Under sustained load, cracks will propagate in hydrogen at

stress intensities greater than a threshold or KXH*

- 132 -
 Temperature C onversions
Albert Sau\eur type i>f table. Luuk up reading in middle column: column ; if in degrees Fahrenheit, read C rntigrade equivalent ii
3 ^
if in degree s Centigrade, read Fahrenheit equivalent in right-hand

- 459.4 t o 0
left-hand column. Valuef^ as printed in "Bethlehem Allo\ SteeU**
s
c ,
Ota 100 100 to 1000 1000 t o 2000 2 0 0 0 t o 3000
3
M

-273
° 1
' -459.4 '
F

-17.8 ' 0
P

32
° 1.
10.0 1
1 ^
50 122.0
C ,
38 too 212
P C
260 500
1 -
932
C

538 2000
H °i
1832 816
P
2500 ' 2732 1 1093
0
2000
P
3632
C
1371 2500
P
4532
-268 -450 1 -17.2 /, 33.8 10.6 St 123.8 43 tlO 230 266 520 990 543 2020 18S0 821 2520 2750 1099 2010 3650 1377 2520 45S0
-262 ' -440 -16.7 2 35.6 11.1 52 125.6 *» 120 248 271 520 1 968 549 1020 1868 827 2520 2768 1104 2020 3668 1382 2520 4568
•
-257 -430 16.1 3 i 37.4 11.7 53 127.4 54 130 286 277 530 986 554 2030 1886 832 2530 2786 1110 2030 3686 1388 2530 4586
-261
-246
-420
' -410
-15.6
-15.0
4
5
39.2
41.0
123
12.8
54 129.2
55 131.0
60 140 284
66 ' J50 ' 302
282
288
540
550
1 1004
1022
560
566
1040 1904
2050 ' 1922
838
843
2540
2550
2804
2822
1116
1121
2040
2050
3704
' 3722
1393
1399
2540
2550
4604
4622
8 oQ
-240 -400 -14.4 6 42.8 13.3 56 132.8 71 160 320 293 560 , 1040 571 2060 1 9 4 0 849 2560 2840 1127 2060 3740 1404 2560 4640 2
-234 -390 -13.9 7 44.6 13.9 1 57 134.6 77 no 338 299 570 1058 577 ! 2070 1958 854 2570 2858 1132 2070 1 3758 1410 2570 4658 H ^
- 229 -3S0 -13.3 S 46.4 14.4 . 5« 136.4 82 ISO ' 356 304 5«0 1076 582 . 2080 1976 860 25S0 2876 1138 2080 ' 3776 1416 2580 4676 tn
-223 -370 -12.8 9 48.2 15.0 59 13B.2 88 190 • 374 310 590 1 1094 588 1 2090 1994 866 ! 2590 ' 2894 1143 2090 , 3794 1421 2590 4694 IS
218 -360 -12.2 10 50.0 15.6 60 1140.0 93 200 392 316 600 1 1112 593 2200 2012 871 1 2600 \ 2 9 1 2 1149 2200 3812 1427 2600 4712 S
CO

-201
-196
212
- 207
-350
-340

-320
330
-11.7
-11.1
- 10.6
-10.0
11
12
13
14
51.8
53.6
55.4
57.2
16.1
16.7
17.2
17.8
61 1141.8
62 143.6
63 1 145.4
64 147.2
99 220 410
100 , 212 413.6
104 220 428
110 230 1 446
321
327
332
338
620
620
630
640
1 1130
1148
1166
1184
599
604
610
616
ttlO
2220
2230
2240
2030
2048
2066
2084
877 ! 2620 ' 2930
882
888
893
2620 , 2948
2630 I 2966
2640 1 2984
1154
1160
1166
1171
2220
2220
2230
2240
3830
3848
3866
3884
1432
1438
1443
1449
2620
2620
2630
2640
4730
4748
4766
4784
i M

H
15 59.0 65 , 149.0 116 240 1 464 650 1202 1250 2102
-190
-184
-310
-300
-
-
9.4
8.9 16 60.8
18.3 ;
18.9: 66 1 150.8 121 250 482
343
349 660 1220
621
627 2260 2120
899 ' 2650 i 3002
904 ' 2660 3020
1177
1182
2250 ' 3902
2260 3920
1454
1460
2650
2660
4802
4820 £:
179 -290 - 8.3 17 62.6 19.4; 67 152.6 127 260 50O 354 670 1238 632 2270 2138 910 2670 3038 1188 2270 3938 1466 2670 4838 2M
- 173
-169
-2)0
273 -459.4
- 7.8
7.2
IS
19
20
64.4
66.2
20.0
20.6'
68.0 21.1
6« 154.4
69 . 156.2
70 ' 158.0
132 ! 270 ' 518
138 280 536
143 290 ! 554
360
366
630
690
700
1256
1274
1292
638
643
649
2280
2290
2200
. 2156
2174
2192
916 2680 1 3056
921
927
2690 3074
2700 ' 3092
1193
1199
2280
2290
3956
3974
2200 ' 3992
1471
1477
1482
2680
2690
4856
4874
4892
c
CO
-168 -270 -454 - 6.7 371 1204 2700
-162 -260 -436 - 6.1 21 69.8 21.7 71 159.8 149 300 . 572 377 720 1310 654 2220 2210 932 2720 3 1 1 0 1210 2220 4 0 1 0 1488 2720 4910
-157 -250 -418 - 5.6 22 71.6 22.2 72 161.6 154 320 . 590 382 720 1328 660 2220 1 2228 938 2720 1 3 1 2 8 1216 2220 4028 1493 2720 4928
-151 -240 -400 - 5.0 23 73.4 22.8 73 , 163.4 160 320 608 388 730 1346 666 1230 2246 943 2730 3 1 4 6 1221 2230 4046 1499 2730 4946
-146 -230 -382 - 4.4 24 75.2 23.3 74 165.2 166 330 626 393 740 1364 671 2240 2264 949 I 2740 3164 1227 2240 4064 1604 2740 4964

-140 -220 -364 - 3.9 25 77.0 23.9 75 167.0 171 340 [ 644 399 750 1382 677 2250 2282 964 ! 2750 3182 1232 2250 4082 1510 2750 4982
-134 -210 -346 - 3.3 26 78.8 24.4 76 168.8 177 350 662 404 760 1400 682 2260 1 2300 960 1 2760 3200 1238 2260 4100 1516 2760 5000
-129 -200 -328 - 2.8 27 80.6 25.0 77 1 170.6 182 360 680 410 770 1418 688 2270 1 2318 966 , 2770 3218 1243 2270 4118 1521 2770 5018
-123 -190 -310 - 2.2 2S 82.4 25.6 78 ; 172.4 188 370 698 416 780 1436 693 2280 2336 971 2780 ' 3236 1249 2280 4136 1527 2780 5036
-118 -ISO -292 - 1.7 29 84.2 26.1 79 i174.2 193 380 716 421 790 1454 699 2290 2354 977 2790 3254 1254 2290 4154 1532 2790 5054

-112 -170 -274 - 1.1 30 86.0 26.7 SO 1176.0 199 390 734 427 800 1472 704 2300 2372 982 1800 3272 1260 2300 4172 1538 2800 5072
-107 -ISO -256 - 0.6 31 87.8 27.2 SI I 177.8 204 400 752 432 810 1490 710 2320 2390 988 2820 ' 3290 1266 2320 4190 1543 2S10 5090
-101 -150 238 0.0 32 89 6 27.8 82 ' 179.6 210 410 770 438 820 1508 716 2320 1 2408 993 2820 3308 1271 2320 4208 1549 2820 5108
- 96 -140 -220 0.6 33 91.4 28.3 S3 181.4 216 420 • 788 443 830 1526 721 1330 1 2426 999 1 2830 3326 1277 2330 4226 1654 2830 5126
- 90 -130 -202 1.1 34 93.2 28.9 M 183.2 221 430 806 449 840 1544 727 2340 i 2444 1004 1 2840 1 3344 1282 2340 4244 1560 2840 5144

- 84 -120 -184 1.7 35 95.0 29.4 85 185.0 227 440 824 454 850 1562 732 2350 2462 1010 1 2850 , 3362 1288 2350 4262 1566 2850 5162
- 70 -110 - 166 2.2 36 96.8 30.0 86 186.8 232 450 842 460 860 1 1580 738 2360 ' 2480 1016 2860 3380 1293 2360 4280 1571 2860 5180
- 73 -100 -148 2.8 37 98.6 30.6 87 188.6 238 460 860 466 870 1598 743 2370 2498 1021 2870 3398 1299 2370 4298 1677 2870 5198
- 68 - 90 -130 3.3 3S 100.4 31.1 SS 190.4 243 470 878 471 S80 1616 749 2380 1 2516 1027 1880 3 4 1 6 1304 2380 1 4 3 1 6 1582 2880 5216
- 62 - SO -112 3.9 39 1023 31.7 S9 192.2 249 1 4S0 896 477 890 1634 754 1390 2534 1032 1890 3434 1310 2390 4334 1688 2890 6234

- 57 - 70 - 94 4.4 40 104.0 32.2 90 ' 194.0 254 3 490 ' 914 482 900 1652 760 , 2400 1 2552 1038 2900 3462 1316 2400 4352 1593 2900 5252
— 51 - 60 - 76 5.0 41 105.8 32.8 91 195.8 488 920 1670 766 2420 ' 2570 1043 2920 3470 1321 2420 4370 1599 2910 5270
- 46 - 50 - 58 S.6 42 107 6 33.3 92 197.6 493 920 1688 771 ! 2420 ' 2588 1049 2920 3488 1327 2420 ' 4388 1604 2920 5288
— 40 - 40 - 40 6.1 43 109.4 33.9 93 199.4 499 930 1706 777 2430 2606 1064 2930 3506 1332 2430 4406 1610 2930 5306
- 34 - 30 - 22 6.7 44 111.2 34.4 94 201.2 504 940 1724 782 1 4 4 0 2624 1060 2940 3524 1338 2440 4424 1616 2940 6324

- 29 - 20 — 4 7.2 45 113.0 35.0 95 ' 203.0 510 950 1742 7 8 8 1 2450 2642 1066 2950 3542 1343 2450 4442 1621 2950 5342
- 23 - 10 14 7.8 46 114.8 35.6 96 204.8 516 960 1760 793 ' 1 4 6 0 2660 1071 , 2960 3560 1349 2460 , 4460 1627 2960 5360
- 17.8 0 32 8.3 47 116.6 36.1 97 206.6 521 970 1778 7 9 9 2470 2678 1077 2970 3578 1354 2470 4478 1632 2970 5378
8.9 4« 118.4 36.7 9S 208.4 527 980 1796 8 0 4 2480 2696 1082 2980 3596 1360 2480 4496 1638 2980 5396
9.4 49 120.2 37.2 99 210.2 532 990 1814 810 2490 2 7 1 4 1088 1990 3 6 1 4 1366 2490 4 5 1 4 1643 2990 5414
37.8 iOO 212.0 538 1000 1832 1093 2000 3632 1649 3000 5432

Reproduced from METAL PROGRESS DATA SHEET No. 18 (1954) by permission

from American Society for Metals, Metals Park, Ohio 44073.
English / Metric (SI) Stress Conversion Factors
Look up stress to be converted in the boldface column If in ksi (103 psi), read MPa in right
hand column If in MPa read ksi in lefthand column Conversion factors 1 MPa 1 MN/m^
(meganev^ton per square metre) or 1 N m m ' (newton per square millimetre) 1 MPa 0 1450377
ksi and 1 ksi 6 894759 MPa

0 to 100 100 t() 200

ksi MPa ksi MPa ksi MPa ksi MPa

0 7 25 50 344 7 14 50 100 689 5 2176 150 1034

0 15 1 6 89 7 40 51 3516 14 65 101 696 4 2190 151 1041
0 29 2 13 79 7 54 52 358 5 14 79 102 703 3 22 05 152 1048
0 44 3 20 68 7 69 53 365 4 14 94 103 710 2 22 19 153 1054
0 58 4 27 57 7 83 54 372 3 15 08 104 717 1 22 34 154 1062
0 73 5 34 47 7 98 55 379 2 15 23 105 724 0 22 48 155 1069
0 87 6 4137 8 12 56 386 1 15 37 106 730 8 22 63 156 1076
102 7 48 26 8 27 57 393 0 15 52 107 737 7 22 77 157 1082
1 16 8 55 16 841 58 399 9 15 66 108 744 6 22 92 158 1089
131 9 62 05 8 56 59 406 8 15 81 109 7515 23 06 159 1096
145 10 68 95 8 70 60 413 7 15 95 110 758 4 23 21 160 1103
160 11 75 84 8 85 61 420 6 16 10 111 765 3 23 35 161 1110
1 74 12 82 74 8 99 62 427 5 16 24 112 772 2 23 50 162 1117
189 13 89 63 9 14 63 434 4 16 39 113 779 1 23 64 163 1 124
2 03 14 96 53 9 28 64 4413 16 53 114 786 0 23 79 164 1131
2 18 15 103 4 9 43 65 448 2 16 68 lis 792 9 23 93 165 1138
2 32 16 1103 9 57 66 455 1 16 82 116 799 8 24 08 166 1145
2 47 17 1172 9 72 67 462 0 16 97 117 806 7 24 22 167 1 151
2 61 18 124 1 9 86 68 468 8 1711 118 813 6 24 37 168 1 158
2 76 19 1310 10 01 69 475 7 17 26 119 820 5 24 51 169 1165
2 90 20 137 9 10 15 70 482 6 17 40 120 827 4 24 66 170 1172
3 05 21 144 8 10 30 71 489 5 17 55 121 834 3 24 80 171 1 179
3 19 22 1517 10 44 72 496 4 17 69 122 8412 24 95 172 1186
3 34 23 158 6 10 59 73 503 3 17 84 123 848 1 25 09 173 1193
3 48 24 165 5 10 73 74 510 2 17 98 124 855 0 25 24 174 1200
3 63 25 172 4 10 88 75 517 1 18 13 125 8618 25 38 175 1207
3 77 26 179 3 1102 76 524 0 18 27 126 868 7 25 53 176 1213
3 92 27 186 2 11 17 77 530 9 18 42 127 875 6 25 67 177 1220
4 06 28 193 1 1131 78 537 8 18 56 128 882 5 25 82 178 1227
4 21 29 199 9 1146 79 544 7 18 71 129 889 4 25 96 179 1234
4 35 30 206 8 1160 80 5516 18 85 130 896 3 26 11 180 1241
4 50 31 213 7 1175 81 558 5 19 00 131 903 2 26 25 181 1248
4 64 32 220 6 1189 82 565 4 19 14 132 910 1 26 40 182 1255
4 79 33 227 5 12 04 83 572 3 19 29 133 917 0 26 54 183 1262
4 93 34 234 4 12 18 84 579 2 19 44 134 923 9 26 69 184 1269
5 08 35 2413 12 33 85 586 1 19 58 135 930 8 26 83 185 1276
5 22 36 248 2 12 47 86 593 0 19 73 136 937 7 26 98 186 1282
5 37 37 255 1 12 62 87 599 8 19 87 137 944 6 27 12 187 1289
5 51 38 262 0 12 76 88 606 7 20 02 138 9515 27 27 188 1296
5 66 39 268 9 12 91 89 613 6 2016 139 958 4 27 41 189 1303
5 80 40 275 8 13 05 90 620 5 20 31 140 965 3 27 56 190 1310
5 95 41 282 7 13 20 91 627 4 20 45 141 972 2 27 70 191 1317
6 09 42 289 6 13 34 92 634 3 20 60 142 979 1 27 85 192 1324
6 24 43 296 5 13 49 93 6412 20 74 143 986 0 27 99 193 1331
6 38 44 303 4 13 63 94 6481 20 89 144 992 9 2814 194 1338
6 53 45 310 3 13 78 95 655 0 2103 145 999 7 28 28 195 1344
6 67 46 3172 13 92 96 6619 2118 146 1007 28 43 196 1351
6 82 47 324 1 14 07 97 668 8 2132 147 1014 28 57 197 1358
6 96 48 3310 14 21 98 675 7 2147 14« 1020 28 72 198 1365
711 49 337 8 14 36 99 682 6 2161 149 1027 28 86 199 1372
7 25 50 344 71 14 50 100 689 5 2176 150 1034 29 01 200 1379

Reproduced from the Mid-June 1980 METAL PROGRESS DATABOOK

by permission from American Society for Metals, Metals
Park, Ohio 44073.
- 134 -
 ENGLISH/METRIC ( S l ) STRESS CONVERSION FACTORS ( C o n t i n u e d )
To convert ksi or MPa values above 400 use the supplementdl table Example Convert 1320 MPa
to ksi Solution 1000 MPa 145 04 ksi (from the small table) and 320 MPa 46 41 ksi (from
the large table) Then. 145 04 46 41 19145 ksi

200 t 3 300 300 t 0 400

ksi MPa ksi MPa ksi MPa ksi MPa

29 01 200 1379 36 26 250 1724 43 51 300 2 068 50 76 350 2 413

29 15 201 1386 36 40 251 1731 43 66 301 2 075 50 91 351 2 420
29 30 202 1393 36 55 252 1737 43 80 302 2 082 5105 352 2 427
29 44 203 1400 36 69 253 1 744 43 95 303 2 089 5120 353 2 434
29 59 204 1407 36 84 254 1751 44 09 304 2 096 5134 354 2 441
29 73 205 1413 36 98 255 1758 44 24 305 2 103 5149 355 2 448
29 88 206 1420 37 13 256 1765 44 38 306 2 110 5163 356 2 455
30 02 207 1427 37 27 257 1772 44 53 307 2 117 5178 357 2 461
30 17 208 1434 37 42 258 1 779 44 67 308 2 124 5192 358 2 468
30 31 209 1441 37 56 259 1786 44 82 309 2 130 52 07 359 2 475
30 46 210 1448 37 71 260 1793 44 96 310 2 137 52 21 360 2 482
30 60 211 1455 37 85 261 1800 45 11 311 2 144 52 36 361 2 489
30 75 212 1462 38 00 262 1806 45 25 312 2 151 52 50 362 2 496
30 89 213 1 169 3814 263 1813 45 40 313 2 158 52 65 363 2 503
3104 214 1475 38 29 264 1820 45 54 314 2 165 52 79 364 2 510
3118 215 1482 38 43 265 1827 45 69 315 2 172 52 94 365 2 517
3133 216 1489 38 58 266 1834 45 83 316 2 179 53 08 366 2 523
3147 217 1496 38 73 267 1841 45 98 317 2 186 53 23 367 2 530
3162 218 1503 38 87 268 1848 46 12 318 2 193 53 37 368 2 537
31 76 219 1510 39 02 269 1855 46 27 319 2 199 53 52 369 2 544
3191 220 1517 3916 270 1862 46 41 320 2 206 53 66 370 2 551
32 05 221 1524 39 31 271 1868 46 56 321 2 213 53 81 371 2 558
32 20 222 1531 39 45 272 1875 46 70 322 2 220 53 95 372 2 565
32 34 223 1538 39 60 273 1882 46 85 323 2 227 54 10 373 2 572
32 49 224 1544 39 74 274 1889 46 99 324 2 234 54 24 374 2 579
32 63 225 1551 39 89 275 1896 47 14 325 2 241 54 39 375 2 585
32 78 226 1558 40 03 276 1903 47 28 326 2 248 54 53 376 2 592
32 92 227 1565 40 18 277 1910 47 43 327 2 255 54 68 377 2 599
33 07 228 1572 40 32 278 1917 47 57 328 2 261 54 82 378 2 606
33 21 229 1579 40 47 279 1924 47 72 329 2 268 54 97 379 2 613
33 36 230 1586 40 61 280 1931 47 86 330 2 275 55 11 380 2 620
33 50 231 1593 40 76 281 1937 48 01 331 2 282 55 26 381 2 627
33 65 232 1600 40 90 282 1944 4815 332 2 289 55 40 382 2 634
33 79 233 1606 4105 283 1951 48 30 333 2 296 55 55 383 2,641
33 94 234 1613 4119 284 1958 48 44 334 2 303 55 69 384 2 648
34 08 235 1620 4134 285 1965 48 59 335 2 310 55 84 385 2 654
34 23 236 162/ 4148 286 1972 48 73 336 2 317 55 98 386 2 661
500 lo 5000
34 37 237 1634 4163 287 1979 48 88 337 2 324 5613 387 2 668
34 52 238 288 2 330 56 27 388 ksi MPa
1641 4177 1986 49 02 338 2 675
34 66 239 1648 4192 289 1993 4917 339 2 337 56 42 389 2 682 72 52 500 344/
34 81 240 1655 42 06 290 1999 49 31 340 2 344 56 56 390 2 689 87 02 600 4137
34 95 241 1662 42 21 291 2 006 49 46 341 2 351 56 71 391 2 696 10153 700 4826
35 10 242 1 669 42 35 292 2 013 49 60 342 2 358 56 85 392 2 703
35 24 243 1675 42 50 293 2 020 49 75 343 2 365 57 00 393 2 710 116 03 800 5516
35 39 244 1682 42 64 294 2 027 49 89 344 2 372 57 14 394 2 717 130 53 900 6205
35 53 245 1689 42 79 295 2 034 50 04 345 2 379 57 29 395 2 723 145 04 1000 6895
35 68 246 1696 42 93 296 2 041 50 18 346 2 386 57 43 396 2 730 290 08 2000 13 790
35 82 247 1 703 43 07 297 2 048 50 33 347 2 392 57 58 397 2 737
2 399 435 11 3000 20 684
35 97 248 1 710 43 22 298 2 055 50 47 348 57 72 398 2 744
36 11 249 1 717 43 37 299 2 062 50 62 349 2 406 57 87 399 2 751 580 15 4000 2/579
36 26 250 1 724 43 51 300 2 068 50 76 350 2 413 58 02 400 2 758 725 19 5000 34 474

Source Anton deS Brasunas University of Missouri Rolla

- 135 -
 6
o
u
C/3 m
>-l CI
a> 0 5 C35 C D OO OO OO OO OO
o , CM CO ^ ^ o
O
.^_> Ol CD o ^— CNJ in
CO
CD
in C O r^ o
•<1- CM C O
•^ in
• *
CO
in CD r^ OO CD
r^ OO CD o
o
CO
CM CO •V in CO r-- CO CD
in CO r^ OO <D o
•^ CM
o CM CO rr rr in CO r- OO
rr in CO r-^ CO CD o CM CO •H
CVJ CO
•^ C O CDCD r^ '— '— • ^ • ^ • ^
'— '— '— CM CM CM CM CM CNJ CM CM CM CO CO CO CO CO CO CO CO CO -a- rr rr
• ^ 03
CO
•H
OO OO OO OO (X3 fs-. fv^ r>_ r._
OO CD CM CO in CD OO CD O CM CO rr rr in C D r-- CO C D o CM CO rr in CM
CO r^
e
a oo CO <T> o CM C O •>r in CO o
C O •=3- in CO OO CD r^
o
CO •O- in CD •^ OO
r^ CD
r^
o
CM CM
CM in CO r--- OO
•^ CD o CM
CM CM CM CM CM CM CM CO CO CO
CO rr CD r^ OO CD
CO CO CO CO CO CO
o CO
rr rr rr rr
(U
p<
o '- -- "-
>.
*i/3 ja
<-i r^ r^ r^ r-^ r^ CO CO CO CO
CO CD O CM CO •"J- in CO OO CD O CM CO rr in CO
h»" r>- • « •
r-^ CO rv. CO CD o OO CM CO • » in CD
1^
r^ OO C-J
O) o CM CO rr in o CM •o- in CO r^ OO CD o CM CD r^ OO CD o
• ^ CM CO in r^
-^ CD o CM CO
> •"3- in CD h - OO CD
"- • ^
CM CM
'- '- CM CM CM CM CM CM CM CO CO CO CO CO CO CO CO CO rr rr rr rr o
o
o % •
co
tD CO CD C O in in in in
in CD r^ co CD O CM CO in• * CO r^ 0 0 CD o CM CO rr in CO r^ OO CD o CM CO rJ- in H CO
t-~
<^
U(/)
CO CO r-
^— CM
O O C3n o CNJ C O
C O in CO I^ OO •^
<j>
o 1 —
CM C O
1 — 1 —
•
^—
CO ^
1 —
r^
^—.»—
O O C D c^
1— CM CM
CM C O <r in r^ C O C D o
CM CM CM CM CM CM CM CO CO CO
CM C O rr in C D O O C D o
CO CO CO CO CO CO
CM C O
rr 5- rr ^ p
M «3-
o
< •
CO
(1> - W
C m ic:
in u-> ID in in
" ^ CO r^ O O C D o
•«r •<3-
CM C O C D
CM C O •<i- CO 1-^ O O C D
•^ •^ •^ in CD l-~ OO CD o CM CO
CVJ CO •<1- in f^ OO CD
,_
o
CM CM
•<3- in CD r^ O O CD C D
CM CO rr in CD OO CD o
,_
CM CO rr
CM
CM CM CM CM CM CM CM CO CO CO
in CO r- OO <D o
CO rr in CO f^ CD
CO CO C O CO CO CO
,_ CM CM
o
CO -a-
CO
rr rr rr rr
§ J=O
^ F T — i ~ • ' — T —
(30 > o
p^ o
M
3 s.'»• -a- •=r •^ •<r C O C O C O CO
^ ,
PH
M
s in C D r^ OO CD o y— CM CO in CD r^ OO CD O o ^ CM
o ^ •>r ir> C O r-- OO o> CD
1^ .)—
CO
CM O CM CO tj- in CO OO CD O
• ^
CM CO in CD l-~ CD o
CM CO
CM
• in CO r^ OO CD
^
CO rr in CD h~ OO o 1 — CM
CO
CO
^ U
CO
is
• ^
CM co •^ in OO CD f~ CM CM CM CM CM CM CM CM CM CO CO CO CO CO CO CO CO CO rr rr rr
• • • • • —
^— • * T— — • • —
'— '— •^ H PH
cs
Wi
CO CO C O C O C O CM CM CM CM
, .. ^ '~^^
C 5 ay
g iHCO
3 CO CO •o- IT) CO r^ OO C D O
CM
O
CO •3- in CO t— OO CD o
CM CO in CO r^ CD
•^ o
CM CO UD CO I-- OO CD o
• ^
CM CO "0- in co r^ OO o
CM
CM
CO ra- in CD r--- OO CD o
CO in CD r^ OO CT>
• ^

CM
CM CO X3 .Q O «
CVJ CO •tr in CO O O CD
"- CM CM CM CM CM CM CM CVJ CM CO C O CO CO CO CO CO CO CO CO rr rr rr 00 4J
^~ ••-
s ^. >*
• ^ • ^
Q.
Ol 0)
cd ^ o. I-H
o
S
(-1
CM C\J CM CM CM
3 "5 (U M
PH CM C\J C O •«r in CO OO
•^ '- cn • ^ •
O
^ y ^
o
CM CO
CM CO
^
in C D r^ CO C D
ir^ CD r^ OO
o
o
,_ CM C O in CD r>. CO CD o
• *
CM CO Tl- in C O r--. CO CD CM
, CVJ CO rr in CD r-. OO CD o
CO rr in CO r^ OO CD o CM CO
, t/>
CO
E E
1 - a rH CO
/^"^ CM CO in C O (^ r- CD •^ CM CM CM CM CM CM CM CM CM CVJ CO CO CO CO CO CO CO CO CO rr rr rr 3
•^ '— ^^ '— •s F IE
>s
• ^
^iH •-1 «d
1 U
C/0 <£ vi
^ ./
•r-
^ ^ ., ^ ^
.»— CM CM
CO in
C D CD CD
CO r- ao CD
^::^ o
o ^— CM
CM C O
CO rr in CO f^ OO CD O
in CO i-~. OO CD
CM CO a- in CD r^ OO CD o
CM CO •<J- r
in CD r^ OO CD o CM
CM CO r3- in CO r~- CO CD o
CO rr in CO r-- OO CD o CO
o
s
o
jac
o S
o
• * • *
C O rl-
T — in CO r^ OO T — T— 1 — T — r- r- T— CM T — CM CM CM CM CM CM CM CM CO CO CO CO CO CO co CO CO rr rr '*"' U
• » — • » —
5 E 1? (U o
"C
^—> o •v^ Xi t4-(
CO
> r u
^
^-A
C3
1
CD
T
o o
CM CO
CM CO

O
•«r in
•

CD CD C D CD C D
in
•>r CO
^

CO
r^ OO
OO CD
r^
o CM cn •<3- in CO 1^ OO C D
CM CO TT in CO 1^ OO C D O
CM
O CM CO rr in C D r--- OO CD
CM CO rj- in C D r-- OO CD o
CM CM CM CM CM CM CNJ CM CO CO
CD
CO
CM CO rr in CO
in CD r^ OO CD
5
CO CO CO CO CO CO CO
r^
o
rr
OO CD CD
CM rr
5- rr rr
o
CO
V
s CP»
E B i g 4J
o 0)
>v
^\ c: o o o.
>
o u •H
<4-lO
* o
O
-«—4
^
>4
^ ^ ^ ^ 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 «2 o o o o

Qi
s
i
o o
o CJ
t/i
•n
(
o
U
W
c
o
CNJ C O
^ m CO r— O O O ) o
O
T— CSI CO
^ m (O T—
1 ^ O O o> o Ol CO ^ CM
m CO t- eo at o CM CO V m CO o &. ^> o £
Oi 3 rt
*s t:O CO 2
c 1^ CM CM CM CM CM CM C«l CVJ CM CO C O C O CO eo CO CO C CO
* • t— H-
•O o
o •H
PJ u u
o. (U
51
 English/Metric (SI) Impact Energy Conversion Factors
Look up impact energy to be converted in boldface column. If in ft-lb, Note. Values for conversions are rounded off to simplify calculations,
read )Oules (J) in lefthand column. If in loules (J), read ft-lb in righthand Actual conversion factors (1 ft-lb = 1.355818 joules and 1 )oule =
column Decimal energy values above 10 and all values above 130 can be 0.737562 ft-lb) can be used if more accuracy is desired. (Reviewed February
calculated by addition. For example, 25.5 ft-lb = 33.9 (25.0)-!- 0.68 (0 5) 1980)
or 34 58 J Also, 165 J = 73 8 (100.0) -h 47 9 (65) or 121 7 ft-lb

0.1 to 20 21 to 130
ioule(J) ft-lb joule (J) ft-lb joule (J) ft-lb joule (J) ft-lb
0.14 0.1 0.07 7.59 5.6 4.13 28.5 zija 15.5 103.0 76Ja 56.1
0.27 0.2 0.15 7.73 5.7 4.20 29.8 22.0 16.2 104.4 77.0 56.8
0.41 0.3 0.22 7.86 5.8 4.28 31.2 23.0 17.0 105.8 nja 57.5
0.54 0.4 0.30 8.00 5.9 4.35 32.5 24.0 17.7 107.1 79.0 58.3
0.68 0.5 0.37 8.13 6.0 4.43 33.9 25.0 18.4 108.4 80.0 59.0
0.81 0.6 0.44 8.27 6.1 4.50 35.3 26.0 19.1 109.8 81.0 59./
0.95 0.7 0.52 8.41 6.2 4.57 36.6 27.0 19.9 111.2 82.0 60.5
1.08 0.8 0.59 8.54 6.3 4.65 38.0 28.0 20.7 112.6 83JI 61.2
1.22 0.9 0.66 8.68 6.4 4.72 39.3 29.0 21.4 113.9 84.0 62.0
1.36 1.0 0.74 8.81 63 4.79 40.7 30.0 22.1 115.2 85.0 62.7
1.49 1.1 0.81 8.95 6.6 4.87 42.0 31.0 22.9 116.6 86.0 63.4
1.63 U 0.89 9.08 6.7 4.94 43.4 32.0 23.6 118.0 87.0 64.2
1.76 U 0.96 9.22 6.8 5.02 44.7 33.0 24.3 119.3 88.0 64.9
1.90 U 1.03 9.36 6.9 5.09 46.1 34.0 25.1 120.7 89.0 65.6
2.03 1.5 1.11 9.49 7.0 5.16 47.5 35.0 25.8 122.0 90.0 664
2.17 1.6 1.18 9.63 7.1 5.24 48.8 36.0 26.6 123.4 91.0 67.1
2.30 1.7 1.25 9.76 7.2 5.31 50.2 37.0 27.3 124.7 92.0 67.9
2.44 1.8 1.33 9.90 7.3 5.38 51.5 38.0 28.0 126.1 93.0 68.6
2.58 1.9 1.40 10.0 7A 5.46 52.9 39.0 28.8 127.4 94.0 69.3
2.71 2.0 1.48 10.2 7J 5.53 54.2 40.0 29.5 128.8 95.0 70.1
2.85 ^l 1.55 10.3 7.6 5.61 55.6 41.0 30.2 130.2 96.0 70.8
2.98 2.2 1.62 10.4 7.7 5.68 56.9 42.0 31.0 131.5 97J) 71.5
3.12 23 1.70 10.6 7J 5.75 58.3 43.0 31.7 132.9 98.0 72.3
3.25 ZA 1.77 10.7 7.9 5.83 59.7 44.0 32.5 134.2 99.0 73.0
3.39 2.5 1.84 10.8 8.0 5.90 61.0 45.0 33.2 135.6 100X1 73.8
3.53 ^6 1.92 11.0 8.1 5.97 62.4 Kja 33.9 136.9 101.0 74.5
3.66 2.7 1.99 11.1 8.2 6.05 63.7 47.0 34.7 138.3 102.0 75.2
3.80 2Ji 2.07 11.3 8J 6.12 65.1 48.0 35.4 139.6 103.0 76.0
3.93 2.9 2.14 11.4 8J4 6.20 66.4 49.0 36.1 141.0 104.0 76.7
4.07 3.0 2.21 11.5 8.5 6.27 67.8 ' 50.0 36.9 142.4 105.0 774
4.20 3.1 2.29 11.7 8.6 6.34 69.1 51.0 37.6 143.7 106.0 78.2
4.34 3.2 2.36 11.8 8.7 6.42 70.5 52.0 38.4 145.1 107.0 78.9
4.47 3.3 2.43 11.9 M 6.49 71.9 53.0 39.1 146.4 108.0 79.7
4.61 3.4 2.51 12.1 8.9 6.56 73.2 54.0 39.8 147.8 109.0 804
4.75 3.5 2.58 12.2 9.0 6.64 74.6 55.0 40.6 149.1 110.0 81.1
4.88 3.6 2.66 12.3 9.1 6.71 75.9 56.0 41.3 150.5 111.0 81.9
5.02 3.7 2.73 12.5 9.2 6.79 77.3 57.0 42.0 151.9 112.0 82.6
5.15 3.8 2.80 12.6 9J 6.86 78.6 58.0 42.8 153.2 113.0 83.3
5.29 3.9 2.88 12.7 94 6.93 80.0 59.0 43.5 154.6 114.0 84.1
5.42 4.0 2.95 12.9 9.5 7.01 81.3 60.0 44.3 155.9 115.0 84.8
5.56 4.1 3.02 13.0 9.6 7.08 82.7 61.0 45.0 157.3 116.0 85.6
5.69 4.2 3.10 13.2 9.7 7.15 84.1 62.0 45.7 158.6 117.0 86.3
5.83 4.3 3.17 13.3 9.8 7.23 85.4 63.0 46.5 160.0 118.0 87.0
5.97 44 3.25 13.4 9.9 7.30 86.8 6M 47.2 161.3 119.0 87.8
6.10 4.5 3.32 13.6 10.0 7.38 88.1 65.0 47.9 162.7 120.0 88.5
6.24 4.6 3.39 14.9 IIJI 8.11 89.5 66.0 48.7 164.1 121.0 89.2
6.37 4.7 3.47 16.3 12.0 8.85 90.8 67.0 49.4 165.4 127.0 90.0
6.51 AA 3.54 17.6 13:0 9.59 92.2 uja 50.2 166.8 123.0 90.7
6.64 4.9 3.61 19.0 14.0 10.3 93.6 69.0 50.9 168.1 124.0 91.5
6.78 5.0 3.69 20.3 15.0 11.1 94.9 70.0 51.6 169.5 125.0 92.2
6.91 5.1 3.76 21.7 1S.0 11.8 96.3 71X> 52.4 170.8 126.0 92.9
7.05 52 3.84 23.0 17J0 12.5 97.6 72.0 53.1 172.2 127.0 93.7
7.19 5.3 3.91 24.4 18J> 13.3 99.0 73.0 53.8 173.5 128.0 944
7.32 SA 3.98 25.8 l9Ja 14.0 100.3 nja 54.6 174.9 129.0 95.1
7.46 5.5 4.06 27.1 20J0 14.8 101.6 75J0 55.3 176.3 130.0 95.9

Reproduced from the Mid-June 1980 METAL PROGRESS DATABOOK by permission from
American Society for Metals, Metals Park, Ohio 44073.

- 137 -
APPENDIX C. MECHANICAL TEST SPECIMENS

APPENDIX C-1

Smooth Bar Tensile Test Specimen A

I
1.875 - 0 . 2 5 " R (Typical)

;^'-14 UNC-2A- 0.020 R Min.

'^ (Typicol (Typical)

•251" I
Chamfer 3 0 " ± 1" .249 -|o.50" k
1.00 -*
To Root Die of Thd
(Typical)
2.00

3.750

- 138 -
APPENDIX C-2

Smooth Bar T e n s i l e Test Specimen B

3/8-16 NC/

* Increase d/ameter from center of gage (.188)

to the ends by 0.002 "

- 139 -
APPENDIX C-3

Smooth Bar Tensile Test Specimen C

1/4-20 UNC

1.63"

- 140 -
APPENDIX C-4

Circumferential Notch at Center of Specimen

for Notched Tensile Specimens

0.189

0.132

S 30°

Notch Root
Radius = 0 . 0 0 5 '

- 141 -
APPENDIX C-5

Tensile Tube Specimen

•^-24NE-F

- 142 -
APPENDIX C-6

Single Edge Notched Tensile Specimen

0.213" Dia-2 Hole:

^ ko.157"

- 143 -
 APPENDIX C-6a

D e t a i l of Notch in Single Edge Notched T e n s i l e Specimen

Notes: 1) All demensions ± O.OOl"

2) Notch root radius = 0.005"
3} To prevent excessive hardening in notch area, machine find
0,040" of notch in five cuts (O.OlO" on first cut, O.OlO"
on 2nd cut, O.OlO" on 3rd cut, O.OOS" on 4th cut and
0.005" on 5th cut).

- 144 -
APPENDIX C-7

C-Shaped Fracture Mechanics Specimen

0.1285 Dia.-2 Holes

r

A k 150"

- 145 -
APPENDIX C-7a

D e t a i l of Notch in C-Shaped Fracture Mechanics Specimen

Notes: 1) All demensions ± 0 . 0 0 1 "

2) Notch root rodius = 0.005"

- 146 -
APPENDIX C-8

Impact Specimen: Modified Naval Research Laboratory

Dynamic Tear Specimen

Mochine Slit 0.04"x 0.170"

- 147 -
APPENDIX D. HEAT ANALYSES

These analyses are for the specific heats of steel actually

tested.

DATA SHEET D-1

Heat Analysis Type 304L Stainless Steel

Element Weight Percent

C 0.03
Mn 1.57
P 0.015
S 0.008
Si 0.43
Cr 18.35
Ni 10.29
Mo 0.17
N -
Al -
Ti -
Nb -
Cu —

148 -
DATA SHEET D-2

Heat Analysis Type 330 Stainless Steel

Element Weight Percent

C 0.049
Mn 1.40
P -
S 0.005
Si 1.46
Cr 18.40
Ni 35.00
Mo 0.18
N -
Al -
Ti 0.45
XTK
WD
Cu 0.20

DATA SHEET D-3

Heat Analysis Incoloy® BOOH

Element Weight Percent

C 0.08
Mn 0.84
P -
S 0.002
Si 0.51
Cr 19.19
Ni 34.04
Mo -
N -
Al 0.36
Ti 0.41
Nb -
Cu 0.52

- 149 -
DATA SHEET D-4

Heat Analysis Tenelon*

Element Weight Percent

C
Mn 15.3
P
S
Si 0.53
Cr 17.4
Ni 0.22
Mo
N 0.4-0.6
Al
Ti
Nb
Cu

DATA SHEET D-5

Heat Analysis Nitronic* 40 Stainless Steel

Element Weight Percent

C 0.015
Mn 9.01
P 0.018
S 0.016
Si 0.24
Cr 20.32
Ni 6.71
Mo -
N 0.35
Al -
Ti -
Nb -
Cu —

- 150 -
DATA SHEET D-6

Heat Analysis Nitronic* 50 Stainless Steel

Element Weight Percent

C 0.05
Mn 5.44
P 0.015
S 0.010
Si 0.42
Cr 21.48
Ni 12.36
Mo 2.12
N 0.25
Al -
Ti -
Nb 0.19
Cu -
V 0.2

DATA SHEET D-7

Heat Analysis Type 316 Stainless Steel

Element Weight Percent

C 0.07
Mn 8.08
P 0.015
S 0.023
Si 0.69
Cr 19.57
Ni 5.67
Mo 2.13
N 0.32
Al -
Ti -
Nb -
Cu -

- 151 -
DATA SHEET D-8

Heat Analysis X18-3 Mn Stainless Steel

Element Weight Percent

C 0.067
Mn 12.4
P 0.013
S 0.013
Si 0.43
Cr 18.55
Ni 3.17
Mo -
N 0.33
Al -
Ti -
Nb -
Cu -
B 0.0015

DATA SHEET D-9

Heat Analysis 18-18 Plus*

Element Weight Percent

C 0.11
Mn 17.80
P 0.020
S 0.004
Si 0.56
Cr 17.78
Ni 0.46
Mo 1.09
N 0.45
Al -
Ti -
Nb -
Cu 0.95
Co 0.01

- 152 -
DATA SHEET D-10

Heat Analysis 304N

Element Weight Percent

C 0.06
Mn 1.66
P 0.30
S 0.025
Si 0.19
Cr 18.37
Ni 8.43
Mo 0.10
N 0.250
Al -
Ti -
Nb -
Cu 0.15

DATA SHEET D-11

Heat Analysis Carpenter 20 Cb-3*

Element Weight Percent

C 0.018
Mn 1.60
P 0.028
S 0.007
Si 0.44
Cr 20.60
Ni 34.90
Mo 4.33
N -
Al -
Ti -
Nb 0.39
Cu 0.20

- 153