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Theses and Dissertations

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The creep and temper embrittlement of Hy-100steel.Courtney Jay Hill

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THE CREEP AND TEMPER EMBRITTLEMENT OF HY-100 STEEL

by

Courtney Jay Hill

A Thesis

Presented to the Graduate Committee

of Lehigh University

in Candidacy for the Degree of

Master of Science

in

Metallurgy and Materials Science

Lehigh University

1975

ProQuest Number: EP76047

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THE CREEP AND TEMPER EMBRITTLEMENT OF HY-100 STEEL

by

Courtney Jay Hill

A Thesis

Presented to the Graduate Committee

of Lehigh University

in Candidacy for the Degree of

Master of Science

in

Metallurgy and Materials Science

Lehigh University

1975

This thesis is accepted and approved in partial fulfillment

of the requirements for the degree of Master of Science.

Professor in Charge

Chairman of Departme

11

ACKNOWLEDGMENTS

The author is sincerely indebted to Dr. Alan W. Pense, for

his advice and constructive criticism, to the Western Electric

Company, for their generous financial support, and to his wife,

Teri, for her patience.

111

TABLE OF CONTENTS

LIST OF TABLES vi

LIST OF FIGURES vii

ABSTRACT 1

INTRODUCTION 3

BACKGROUND 3

TEMPER EMBRITTLEMENT LITERATURE REVIEW 5

CREEP EMBRITTLEMENT LITERATURE REVIEW 15

EXPERIMENTAL PROCEDURE 20

INTRODUCTION 20

MATERIAL AND HEAT TREATMENT 21

STEP COOLING 22

ISOTHERMAL AGING 23

CREEP EMBRITTLEMENT 23

MECHANICAL TESTING 24

RESULTS AND DISCUSSION 28

HEAT TREATMENT 28

TEMPER EMBRITTLEMENT 30

CREEP EMBRITTLEMENT 33

OPTICAL AND ELECTRON MICROSCOPY 35

CONCLUSIONS 37

TABLES 39

FIGURES 45

IV

REFERENCES 62

APENDICES A-l

APPENDIX A: CHARPY CURVES - TEMPER EMBRITTLED HY-100 A-l

APPENDIX B: CHARPY CURVES - CREEP EMBRITTLED HY-100 A-21

LIST OF TABLES

Table Title

L Chemical Analysis of HY- 100

2 Effect of Quench Media on Center Ha If-Time-Temperature Cooling Rate

3 Mechanical Properties of HY-100 Tempered at 1100 and 1200°F

4 Effect of Step Cool Embrittlement on Impact Properties

5 Effect of Isothermal Aging on Im- pact Properties

6 De-embrittlement of HY-100 at 1050°F for 1 Hour

7 Mechanical Properties of HY-100

8 Effect of Applied Stress on Em- br itt lement

VI

LIST OF FIGURES

Figure Title

1 Cooling Study - Plate Specifications

2 Preparation of Impact Specimens from Creep Specimen

3 Specimen Orientation

4 Effect of Plate Thickness on Center Cooling Rate

5 Tempering Dependence of Hardness

6 Tempering Dependence of Strength

7 Tempering Parameter

8 Correlation of Mechanical Properties

9A-9B Kinetics of Isothermal Embrittlement - 110 ksi Yield Strength

10A-10B Kinetics of Isothermal Embrittlement - 135 ksi Yield Strength

11A-11B Effect of Embrittlement on the Upper Shelf

12A Optical Micrograph of Quenched and Tempered HY-100

12B SEM Micrograph of Quenched and Tempered HY-100

13 SEM Fractograph - Step Cool Temper Embrittled

14 SEM Fractograph - Step Cool Temper Embrittled - Subsequently De- embrittled

15 SEM Fractographs of Temper Embrittled HY-100

VII

ABSTRACT

The creep and temper embrittlement susceptibility of medium

section (4 inch wall thickness) HY-100 quenched and tempered to

110 and 135 ksi yield strength was examined by simulating heat

treatments with plates 1 1/2 inches thick. Embrittlement was

produced by step cooling, isothermal aging at 750°F, and creep

aging at 750°F. The severity of embrittlement was evaluated by

measuring the shift of the midrange brittle to ductile transition

temperature of Charpy V-notch specimens, as determined by the

absorbed energy and the lateral expansion criteria. Uniaxial

tension tests, performed on unnotched specimens at room tempera-

ture and notched specimens at -58°F, were also employed in an

attempt to assess the influence of embrittlement on tensile

properties. Optical and scanning electron microscopy were uti-

lized to examine the microstructural and fracture characteristics

of creep and temper embrittlement.

The overall impact transition temperature of the steel as

quenched and tempered was quite low on the order of -30°F to

-110°F, however the shelf toughness of the alloys was also low,

about 35-45 ft-lb. Embrittlement of the quenched and tempered

steel was revealed to be severe, on the order of 300°F for step

cooling and 165°F as measured by isothermal aging for 5000 hours.

However, the specific transition temperature change was influenced

1

by Charpy impact specimen size, the impact resistance criterion

employed and the steel strength level. While step cooling

provided an estimate of the maximum extent of embrittlement

possible in this steel, it does not exactly reproduce actual

service at 750°F, thus it probably is an overestimate of embrittle-

ment in service. Step cooling also produced a deterioration in

the shelf toughness that was not observed in the isothermally aged

steel. De-embrittlement at 1050°F for 1 hour of the embrittled

steels demonstrated the reversibility of temper embrittlement.

Creep embrittlement, characterized by a degeneration of

impact properties at ambient temperatures and grain boundary void

formation, was not observed. Rather, a delay of temper embrittle-

ment, particularly at the lower strength level, was recorded.

INTRODUCTION

BACKGROUND

Low a Hoy high strength steels have been used in a variety

of industrial applications involving elevated temperature exposure

for over 50 years. In spite of their generally good service record,

the long time application of these steels for this type of service

has been of increasing concern in recent years since the possibil-

ity that metallurgical aging effects can change the mechanical

properties of the material in service has become a demonstrated

reality. In many instances, these changes are induced by the effects

of temperature alone, but in others, plastic strain and/or environ-

mental influences may be significant. Examples of such effects are

decreased creep ductility or stress corrosion cracking in Cr-Mo

oil refinery equipment, increases in strength and decreases in

ductility in austenitic stainless alloys, and, most commonly,

decreased impact toughness in heat treated steels due to temper

embrittlement. This latter effect has been identified as a sig-

nificant cause for concern in both the electric power generation

and petroleum industries. The American Society for Testing and

Materials, The American Petroleum Institute, and The Electric

Power Research Institute have initiated important research programs

in this area in recent years. Work sponsored by A.P.I, at Lehigh

University established that, during the first year of service at

temperatures of 800 to 900°F, the impact transition temperature

3

of 2 1/4 Cr - 1 Mo steeL pressure vessels (ASTM A542) may increase

as much as 150°F. As a result, the 70°F Charpy impact toughness

may be in the range of 10 ft-lbs or less. This low toughness in

the ambient temperature range has resulted in at least two cases

of major cracking, necessitating expensive repairs in refinery

vessels in both the United States and Japan. In addition, there

is at least one instance of a major failure in power generation

equipment in Great Britain that was, in part, attributable to the

low toughness induced by temper embrittlement. It is prudent,

therefore, that users of low alloy steels which operate in the

elevated temperature range consider the effects of service envi-

ronments on the static ambient and elevated temperature properties

of these steels.

It is the purpose of this research to develop precisely such

information on the effects of a moderate service environment on

the toughness of HY-100 type (Ni-Cr-Mo) steel after an extended

period of service. This type of steel is used as plates (HY-80,

HY-100, ASTM A543), forgings (A508 Class 4), and other product

forms in a number of pressure vessels including nuclear pressure

vessels. The particular application of interest in this program

is service in intermediate thickness (4 inch wall thickness)

forged vessels used to grow quartz single crystals in caustic

solution. The temperature of service in this case is up to ap-

proximately 750°F, and pressures at this temperature are relatively

4

high. Although pressures in the ambient temperature range are

low, the relatively high pressures in service dictate that these

steels, for design purposes, be heat treated to high strength

levels. While, indeed, the service environment is only moderate

with respect to those which produce normal temper and creep

embrittlement effects, the long times of service (greater than

50,000 hours) make such effects a possible source of embrittle-

ment. For this reason, a three phase study of HY-100 Ni-Cr-Mo

steel at 110 and 135 ksi yield strength levels to isothermal

embrittlement was performed. The three phases include step

cooling and isothermal aging and creep testing at 750°F.

Along with this experimental program, a literature review of

this topic was also undertaken. This literature review is pre-

sented in the next sections of the Introduction to this report.

TEMPER EMBRITTLEMENT LITERATURE REVIEW

Although long recognized to be of significant metallurgical

importance — the first published reference appeared in 1917 --

the theory of temper embrittlement has eluded complete develop-

ment. The problem approached major proportions during the con-

certed effort to utilize alloy steels in World War I. In particu-

lar Krupp of Germany produced a Ni-Cr armor steel which, in heavy

(2) sections, was extremely susceptible to brittle failure. ' The

manifestations of temper embrittlement were further demonstrated

in 1954 when several large steam turbine generator rotors, within

(3 4) a period of ten months, experienced catastrophic failure. '

LJ5) In 1969, yet another steam turbine rotor failed catastrophica II3

Each of these failures, to some extent, was attributable to

temper embrittlement. Failure caused by temper embrittlement has

also been experienced in the petro-chemica1 industry and is of

vital concern to operators of any large, heavy section chemical

pressure vessels.

The long history of failures, in spite of extended and exten-

sive metallurgical research, indicates the complex and perplexing

nature of the temper embrittlement phenomenon. In general, it

is a phenomenon that, although well characterized, is not well

understood. The extent of the present understanding can best be

presented by listing the prominent characteristics of temper

embrittlement. This list represents a modification and extension

of that developed by Emmer, Clauser and Low.

Mechanical Properties

1. Temper embrittlement is most commonly character-

ized by a reduction of impact resistance, and is

usually described by the shift in the ductile to

brittle transition temperature of the steel. ' '

2. Embrittled steels are sufficiently tough that an

impact test is usually required to detect the

phenomenon.

3. Shelf energies are not affected by isothermal

embrittlement. However, a step-cooling assess-

6

merit of temper embrittlement may reduce shelf

(10) energies .

4. Embrittlement does not affect those mechanical

properties of the material usually determined

by the tension test. It may, however, affect

flow strength.

5. Temper embrittlement does not affect the

endurance limit of the material, although it

does cause a reduction of fatigue strength

above it. At high stress levels this effect

, (12,13,14) is not pronounced. ' '

6. The rate of fatigue crack propagation increases

(14) with the degree of embrittlement.

7. Materials that are susceptible to temper

embrittlement are also susceptible to hydrogen

embrittlement.

Kinetics

1. Temper embrittlement occurs in the temperature

range of 700-1000°F(?'16'U'L8) and exhibits a

typical C-curve time-temperature relationship

within this range.

2. At low temperatures within the embrittling

range, the initial rate of embrittlement is

slow, while the maximum embrittlement and the

u- • A (19) time to achieve it are increased.

7

3. After an extended period of time, the rate of

K .--i - • A A (7,8,20,21) embrittlement is reduced.

4. Overaging appears to occur at temperatures within

the embrittlement range somewhat above the nose

of the C-curve. However, this apparent overaging

must also take into account the tempering of the

^ • i „t. , (8,16,20-24) material that occurs. ' ' '

Alloy effects

1. The susceptibility of a steel is influenced by

& its alloy content.

2. Plain carbon steels are not observed to

embrittle when subjected to an embrittling

. (7,22,25) treatment. ' ' '

3. Cr, Ni and Mn increase susceptibility to temper

embrittlement. However, specific effects are

determined by the interaction between these and

other alloy or impurity elements present in the

„ . (8,22,25,26) steel. ' ' ' '

4. Mn and Cr individually must be present in a mini-

mum concentration in order to induce temper

embrittlement. This minimum concentration is

influenced by the specific impurity element

(1 22 26^ concentration of the alloy. > ' '

5. Susceptibility of a steel is related to harden-

ability with respect to Ni, Cr and Mn. This

8

effect is amplified by the fact that these

elements tend to strengthen the grain relative

to its boundary, thereby shifting the fracture

transition temperature to higher temperatures.

6. In alloy steels, lowering the carbon content

decreases the degree of embrittlement. This

effect is obscured by the fact that reducing

the carbon level of the steel decreases the

transition temperature in the unembrittled

condition. A complete absence of carbon, how-

(22 25) ever, does not prevent embrittlement. '

7. Embrittlement occurs only in the presence of

specific impurity elements. These are Sb, P, Sn,

and As in order of decreasing severity on a

weight basis. Bi, Se, Ge, V, and Te have also

been demonstrated to promote temper embrittle-

^ (22,25,27-30) ment. ' ' '

8. In addition to promoting temper embrittlement,

the addition of P and Sb to alloy steels causes

an increase in the transition temperature of

the unembrittled material. ' '

9. Controlled amounts of Mo, W, and Ti suppress

both the rate and the extent of temper embrittle-

. (20,28,52) ment.

10. Mo retards both embrittlement and de-embrittlemenr.

11. Mo and W must be present in a minimum con-

centration in order to suppress embrittle-

nient . A deviation from this amount will

reduce their effectiveness and may even

result in increased brittleness.

12. Alloy and impurity interactions have a sub-

stantial effect on temper embritt lenient. V

and Mo alone, for example at a concentration

of 1%, have been shown to induce embrittle-

" ment. However, a 1% Mo - 1% V addition

«..VT. (20,55) decreases susceptibility. '

Segregation

1. Both impurity elements and alloying elements

are concentrated along prior austenite grain

(32-37) boundaries during embrittlement.

However, segregation has been shown to occur

along all high angle grain boundaries of an

embrittled steel. This segregation is

significant when the angle of grain boundary

- *. • «- «-u ic A (19,37,38) mismatch is greater than 15 degrees. ' '

2. Cr, Ni, and Mn segregate to an appreciable

extent only in the presence of minor impurity

elements. '

3. The depth of impurity segregation in an

embrittled steel is on the order of 5 to 10 A.

10

The segregation profile of alloy elements tends

(31) to be more diffuse, approximately 50 A. The

concentration of the segregated species is reported

(31,33,40) to be at least five times the bulk concentration.

4. Segregation does not occur in the austenite field

but in the ferrite plus carbide field. Segregation

in the ferrite phase is also reversible in the

f (19,31) ferrite phase. '

5. There is no immediate correlation between the

amount of segregation and the degree of de-

U ■ 1 ■ U A ■ (19) cohesion along gram boundaries.

Reversibility

1. Temper embrittlement may be reversed by heating

(20 32) above the embrittling temperature range. '

2. The kinetics of the de-embrittlement process

are influenced by alloy composition and temper-

ature. The addition of Mo, for example, sup-

presses the rate of reversibility as well as

(20 32 52) the rate of embrittlement. ' '

3. Higher temperatures permit shorter times for

(32) de-embrittlement.

Surface effects

1. In general, embrittlement is characterized by

a change in fracture mode from trans- to

(8 37^ intercrystalline. ' This does not require

11

an associated large shift in the ductile to brittle

(25) transition temperature of the steel. However,

both the percent intergranular fracture and the

percent grain boundary decohesion, as well as the

numerical shift in transition temperature, have

been employed to assess the extent of embrittlement,

2. The grain boundary facets which characterize the

fracture surfaces of embrittled steels are in-

distinguishable from those observed on the un-

embrittled material. However, their chemistry is

(32) quite different.

3. In low alloy steels, fracture is observed to

occur along prior austenite grain boundaries.

In carbon free alloys fracture proceeds along

(37) ferrite-ferrite grain boundaries.

4. Temper embrittlement may sometimes be deduced in

tensile specimens by a radial crack pattern along

T- c r (2.30) the fracture surface.

5. In steels containing chromium and phosphorous,

an etheral picric acid etch will preferentially

(41) attack prior austenite grain boundaries.

Heat treatment

1. The upper shelf fracture mode of some temper

embrittled steels is sensitive to prior heat

treatment. In some 3.5% Ni - 1.7% Cr based steels,

12

a step austenitizing treatment prior to embrittle-

ment results in a reduction of the upper shelf and

a change from intergranular fracture to intergranu-

lar microvoid coalescence. The transition tempera-

(42) ture, however, is unaffected by this treatment.

2. Prolonged heating at a temperature above the

embrittling range inhibits subsequent temper

(8 30 32) embrittlement by reducing the embrittling rate. ' '

3. Intercritical heat treatment has been observed to

reduce susceptibility to temper embrittlement by

manipulation of the microstructure such that the

interacting alloying and impurity elements are

U 1 • 11 + A (43) morphologically separated.

Microstructural effects

1. Susceptibility to temper embrittlement is in-

fluenced by hardness but these effects are not

entirely clear. Steels possessing higher yield

and tensile strengths, (i.e., in the quenched and

tempered condition rather than a softer one), will

more severely embrittle. However, for steels with

approximately the same microstructure, the higher

strength steel appears to be embrittled less than

the lower strength one.

2. Martensitic steels undergo a larger shift in

transition temperature than do bainitic steels.

13

Bainitic steels, in turn, are more susceptible

than pearlitic steels. The embrittled transition

temperature, however, may be independent of struc-

ture. The different shifts result from differing

'..,.. i u -..I A J-.- (8,26,56) initial unembrlttled conditions. '

3. Increased austenitic grain size results in

A u .^i ,. (7,44,45,46) increased embrittlement.

4. Plastic deformation prior to embrittlement

reduces the susceptibility of quenched and

v tempered alloy steels. Plastic deformation during

embrittlement appears to delay the development of

temper embrittlement. The extent of embrittlement,

however, may be increased by the initiation of

creep embrittlement. Plastic deformation sub-

sequent to embrittlement lowers the transition

* ^u .i (47,48,49) temperature of the steel. ' '

From examination of this list it is understandable that, at

present, no well formulated model for the temper embrittling

process exists. While it is clear that impurity and alloy segre-

gation play a major role in the process, it is not clear to what

extent each of these is responsible or in what manner these species

become segregated in the steel.

While recent results indicate that segregation occurs during

embrittlement in the ferrite phase, and does not proceed by any

mechanism requiring pre-segregation during austenitizing, the

14

question of how these elements concentrate along high angle grain

boundaries remains unanswered. It is currently believed that this

• i -v • ... (20,31,36,50) ■ .' occurs by either equilibrium segregation or by alloy

• «.. J • • ... . • * -1. • u J • (18,24,37,51,52) rejection during precipitation at the grain boundaries. '

The difficulty in justifying both of these mechanisms is that

those characteristics of temper embrittlement which are exploited

to defend one model may also be suitably tailored to support the

other. Specifically, the kinetics, the reversibility, the segre-

gation and alloy effects may be the consequence of either an

equilibrium segregation or a precipitation mechanism.

CREEP EMBRITTLEMENT LITERATURE REVIEW

To present, creep embrittlement has not received the metal-

lurgical attention reserved for temper embrittlement. However,

it, too, has long been recognized to have profound effects upon

ultimate material performance.

Creep embrittlement became recognized as a phenomenon of

consequence when boiler flange bolts manufactured of heat resisting

steel experienced failure within several years of service. '

Specifically, it had been found that vessels operating in a temper-

ature range of 850 to 950°F experienced catastrophic failure at

the bolts within a service period of 3000 to 12,000 hours, depend-

(59) ing upon the composition of the steel and the loading stresses. '

Creep embrittlement must also be considered in those appli-

cations for which temper embrittlement may be significant. For,

in general, those materials which are susceptible to temper embrit-

15

tlement in service are, under particular conditions of temperature

and stress, likely to exhibit creep embrittlement. Hence, manu-

facturers of large pressure vessels and petro-chemical equipment

such as hydrocracking towers and reformers must also consider

creep embrittlement during design and material evaluation.

Presented below are the most prominent characteristics of creep

embrittlement.

Mechanical properties

L. Creep embrittlement is accompanied by the

reduction, and subsequent partial recovery of

stress rupture ductility. '

2. Associated with creep embrittlement is the loss

of notch ductility and an increase in the

brittle to ductile transition temperature of

„. _ . (61,62) the steel.v '

3. A reduction of the upper shelf energy has also

been reported to be a consequence of creep

embrittlement. Hence embrittlement may be

characterized by a reduction in the ambient

temperature impact resistance of the steel.

Kinetics

1. In ferritic steels, creep embrittlement occurs

in the temperature range of 800 to 1100°F. (65_6V3)

2. At lower temperatures within the embrittling

range, the rate of creep embrittlement diminishes.

16

However, susceptibility increases and embrittlement

*«. , (65,70-72) maximizes after longer exposures. '

3. The most severe conditions for a given material

are those which produce a small amount of creep

*- - A A .AC*- (65,67,70,72) after an extended period of time. ' ' '

Alloy effects

1. There is some evidence that the temper embrittle-

ment inducing elements may also contribute to the

creep embrittlement of steels. '

2. In general, those materials susceptible to temper

embrittlement are also likely to exhibit creep

embrittlement under the appropriate service

,._. (60-62) conditions.

Reversibility

1. Creep embrittlement, to distinguish it from temper

embrittlement, is irreversible by subsequent heat

treatment. However, some improvement of the

transition temperature may be produced by tempering

effects. In effect, the void formation associated

f 62 "4 with creep embrittlement is irreversible.

Surface effects

1. Creep embrittlement is accompanied by a transition

from a transgranular to intergranular mode of

(73) fracture.

17

Microstructure effects

1. Creep embrittlement is characterized by the

formation of small voids and microcracks along

prior austenite grain boundaries. These are

commonly observed to be formed transverse to the

., .,- • (62,64) tensile direction.

2. For a tempered martensitic structure, suscepti-

bility is observed to be greater at higher strength

levels. <65>67>

3. Coarse grain materials are more susceptible than

(65,67) fine grain ones. '

The mechanism by which creep embrittlement proceeds is

intimately associated with the specific carbide precipitation

processes and kinetics that occur tinder the specific conditions

of stress and temperature. During exposure to elevated tempera-

tures, carbides are precipitated at austenite grain boundaries.

Simultaneously, alloy carbides are formed within each grain and

tend to dispersively strengthen it. However, it is observed that

those carbides which are precipitated near grain boundaries tend

to dissolve and reprecipitate on the complex grain boundary

carbides. As a result a thin, weak denuded zone is formed. The

combination of a strong grain bulk, a weak grain boundary region,

and a stress concentrating grain boundary precipitate promotes

intergranular void formation under an applied stress which even-

tually leads to failure. Consequently, the susceptibility of a

18

steel to creep embrittlement is determined by the manner and extent

to which the strength of the grain interior is increased by the

dispersion of incoherent carbide precipitates relative to the

weakening of the grain boundary region by denudation and stress

concentrating grain boundary precipitation. When the strength of

the individual grains is enhanced to such an extent that bulk

deformation is preceded by grain boundary sliding and void for-

mation, embrittlement results and the fracture mode shifts from

trans- to intergranular. The precipitation process during creep

embrittlement not only acts to promote intergranular failure but

also prohibits the normal accommodation of strain under creep

conditions, such as bulk deformation, climb, and grain boundary

migration.

( f\fi \ Roper summarizes these individual contributions as

fo Hows :

1. The fine dispersion of carbide precipitates within the

grain interiors produces significant resistance to bulk

deformation of the individual grains during creep.

2. The increasing width of the denuded zone with time at

temperature lowers creep strength near the grain bound-

aries .

3. The formation of grain boundary carbides reduces the

cohesive strength of the boundaries and the combination

of hard particles in a soft matrix, as stress concen-

trators, initiate voids.

19

EXPERIMENTAL PROCEDURE

INTRODUCTION ?

In order to ascertain the susceptibility to temper and

creep embrittLement of heavy section commercial grade HY-100,

it was necessary to simulate the microstructure of a heavy plate

given an industrial heat treatment, using a smaller plate thick-

ness that could be treated in the laboratory. In this case, the

heat treatment of a 4-inch thick water quenched plate was simu-

lated with one 1.5 inches thick to obtain 1) the desired center

half-time-temperature cooling rate upon quenching from the

austenitizing temperature, and 2) the desired yield strength

levels in the tempered condition, 110 and 135 ksi. Heat treated

specimens were subsequently embrittled isothermally at 750°F and

by a standard step cooling process. Standard Charpy impact,

uniaxial room temperature tensile, and low temperature notched

tensile tests were performed on the aged material to assess the

extent of embrittlement. Reversibility of the temper embrittle-

ment was emphasized by de-embrittlement of some isothermally aged

and step cooled specimens.

To investigate the effect of an applied stress on the

embrittlement process, creep specimens were loaded to one half

their quenched and tempered yield strength at 750°F. Subsequently,

Charpy data was accumulated by utilizing friction welded subsize

(5 by 5 by 55 mm) V-notch impact specimens.

20

Finally, to further characterize the microstruetura 1 ramifi-

cations of temper and creep embrittlement in HY-100, standard

metallography and scanning electron metallography were performed.

MATERIAL AND HEAT TREATMENT

The material used in this investigation was HY-100 quenched

and tempered steel in the form of a 23" dia. forging cut into

1 1/2 inch slices. The material was obtained from Earle M. Jor-

gensen Co. and the composition is listed in Table 1. To reproduce

the water quenching of a 4-inch plate steel in a 1.5-inch plate,

a simulation procedure based on cooling rates was employed. Al-

though previously used at Lehigh in several programs, the effect

of various quench media on the cooling rate of 6 in. by 10 in.

by 1.5 in. HY-100 plate was again explored to verify previous

work and to establish the exact procedures necessary to heat

treat the steel in this program. The cooling rates were deter-

mined by inserting three shielded chrome1-alumel thermocouples

into the center thickness of the test plate, austenitizing for

1 hour at 1575°F, and independently recording the subsequent

cooling process with three strip chart recorders (see Figure 1).

Variations of cooling rate, as measured at each couple position,

were essentially negligible. This procedure was repeated for a

series of coolants to determine which would reproduce the 4 inch

cooling rate in the 1.5-inch plate most closely. The results of

this study are discussed later.

21

To obtain the desired strength levels, 110 and 135 ksi

quenched and tempered yield strengths, it was necessary to

identify the tempering dependence of the mechanical properties

of HY-100. This time-temperature relationship was accomplished

by tempering 1 in. by 10 in. by 1.5 in. sections for 1, 2, 5, 10

and 30 hours at 1100°F and performing an identical treatment at

1200°F. Following the temper, Brinell hardness tests, utilizing

a 3000 kg. load and tensile tests, were performed. As is

described later, from these tempering relationships, the appro-

priate heat treating times and temperatures were determined.

This determination was substantiated by mechanical testing of

material receiving these treatments.

STEP COOLING

An immediate assessment of the susceptibility of HY-100 to

temper eifibrittlement wasobtained by utilizing the following

step cooling treatment on 6 in. by 5 in. by 1.5 in. plates at

each strength level:

Holding 1 hour at 1100°F followed immediately by

15 houts at 1000°F " " "

24 hours at 975°F " " "

48 hours at 925°F " " "

72 hours at 875°F " " "

air cooling to room temperature.

Treatment was done in a forced air furnace and temperature

measurements performed during the treatment demonstrated that the 22

time between each cooling step was a small fraction of the total

time at any given temperature as outlined by the required treat-

ment .

ISOTHERMAL AGING

Isothermal aging was accomplished by subjecting 6 in. by 5 in.

by 1.5 in. plates at each strength level to 750°F for 1, 10, 30,

100, 300, 1000, 2000 and 5000 hours in a forced air furnace.

Specimens were immediately water quenched from the aging temper-

ature.

CREEP EMBRITTLEMENT

To evaluate the influence of the addition of stress to the

susceptibility of HY-100 to embrittlement, 0.375 inch diameter,

6 inch gage length creep specimens were subjected to a stress

level equal to one half the yield strength for 1000 hours at

750°F. This procedure was repeated for each strength level.

Subsequently, impact specimens were prepared by cutting

1 inch sections from the gage and friction welding each of these 4

between two 1-inch lengths of 1020 steel bar stock (see Figure 2).

These welded assemblies were then ground and milled into 5 by 5

by 55 mm Charpy V-notbh impact specimens (these are subsize speci-

mens machined according to ASTM E23 specification). This pro-

cedure was followed in order to conserve test material and to

reduce the time necessary to obtain a sufficient number of Charpy

specimens from the crept material to determine a full impact curve

Since only the material in the vicinity of the notch is actually

23

tested, the sectioning and welding procedure remains valid as

long as the welding process does not alter the material near the

notch. For this reason, friction welding, characterized by a

small heat affected zone, was employed to fabricate the assemblies.

This procedure has been used in previous studies at Lehigh

TT •- 60 University.

Because such a small notch (0.39 inch deep) is required

in the subsize Charpy specimens that were used, notch uniformity

is essential. This was verified by comparing the profile of

optically magnified notches, and measuring the width of each

notch at the specimen surface to determine variability.

Since specimen size effects are important in impact tests,

subsize Charpy specimens of the quenched and tempered and of the

step cooled material were also prepared. In this manner, the

impact data accumulated from creep embrittled specimens could be

directly compared to unembrittled and step cool embrittled

material of the same specimen size.

MECHANICAL TESTING

Mechanical test specimens were cut from the center thick-

ness of the test plates, and surface layers were discarded. A

schematic diagram of the orientation of the test specimens is

found in Figure 3.

Tensile Tests

Tensile testing was performed at room temperature on

0.252 inch diameter specimens. All tests were performed at a

24

strain rate of 0.02 in/min. Elongation prior to yielding was

measured with an extensometer. A second set of specimens was

prepared with a 0.05 inch deep, .005 root radius notch at the

center of the gage length. These were tested at -58°F also at

a strain rate of 0.02 in/min. Data obtained from these tests

iflac luded yield load, as determined by a 0.027° offset, tensile

load, fracture load, reduction in area, and percent elongation.

Yield, tensile and fracture loads were subsequently converted to

the appropriate values of engineering stress.

Impact tests

Chapry V-notch impact curves were employed to determine

the effect of embrittlement on the upper shelf energy and the

transition temperature of the quenched and tempered material.

To develop full curves, both high and low temperature testing

was required. Low temperatures were obtained by utilizing

liquid nitrogen to cool a 2-methylbutane bath. Heated water

was used to obtain temperatures from room temperature to 150°F.

Boiling methanol-water solutions were found to produce stable

temperatures from 150 to 212°F. For higher temperatures, heated

oil was employed. All specimens were immersed in the appropriate

bath for a minimum of 5 minutes before testing. Subsequently,

fractured specimens were immediately submerged in acetone to

prevent frost formation on low temperature specimens and to remove

the oil from the very high temperature specimens. Fracture sur-

faces were then dried in air and immediately sprayed with an

25

acrylic coating to protect them from contamination. For each set

of specimens, the characteristic transition temperature and shelf

energies were determined from curves fitted with the aid of a

special computer program developed and on file at Lehigh University.

Metallography and Surface Preparation

To examine the micros truetura1 effects of creep embrittle-

nient on HY-100, samples of the quenched and tempered and creep

embrittled material were given a standard metallographic polish.

Creep specimens for metallographic examination were sectioned

longitudinally prior to examination.

The particular metallographic procedure employed required

the mounting of specimens in bake lite before polishing. Initially,

mounted samples were ground on an 80 grit wet belt sander to

develop a flat surface. To produce a final polish acceptable for

examination at high magnifications, surface preparation was

completed with 6 micron and 1 micron diamond paste and a quick

buff on the Linde B (0.06 micron alumina) polishing wheel. Etching

was performed by swabbing each polished specimen with 2% nital for

approximately fifteen seconds. Typically, the surface was lightly

repolished with Linde B and re-etched. This procedure was repeated

until a satisfactory surface could be obtained.

Examination of the prepared specimens in the optical

microscope was not entirely satisfactory because the fine structure

could not be resolved. Consequently, microstructura1 analysis was

completed with the aid of the scanning electron microscope.

26

Fracture surfaces did not require any special surface

treatment prior to examination with the scanning electron micro-

scope. However, stripping of the protective coating, degreasing,

and cleaning were performed by immersing the fracture surfaces in

.acetone and ultrasonically cleaning them for 15 minutes.

Scanning Electron Microscopy

Scanning electron microscopy was utilized to examine

fcjacture surface appearance of Charpy impact and tensile specimens.

Specifically, each of the following were inspected under the 20 KV

electron beam and photographed on Polaroid PN Type 55 film.

1. Unnotched tensiles - unembrittled and

temper embrittled

2. Notched tensiles - unembrittled and temper

embrittled

3. Upper shelf Charpy - unembrittled and

progressively embrittled

4. Lower shelf Charpy - embrittled

5. Lower shelf Charpy - de-embrittled

6. Upper shelf Charpy - creep embrittled

In addition, microstructura1 information was acquired

from polished and etched specimens of HY-100 in the unembrittled

and creep embrittled conditions.

27

RESULTS AND DISCUSSION

HEAT TREATMENT

An examination of Figure 4 illustrates that to simulate the

cooling characteristics of a water quenched 4 inch plate, a center

cooling rate in a 1.5 inch plate between approximately 1.8 and

3.1°F/sec is required. To produce this rate, several coolants

were tested. Presented in Table 2 are the center thickness half-

time- temperature cooling rates derived from the various quench

media employed. It is evident from these data that the most sat-

isfactory results were achieved by forcing room temperature air

over the plate faces with high capacity rotary fans. During the

course of experimentation, it was determined that while the

desired simulated cooling rate could be approached by placing

one fan perpendicular to each side of the austenitized plate at

a distance of 8 inches from the surface, it was possible to obtain

slightly faster cooling by employing two additional fans.

Utilizing this procedure, a consistent, reproducible value of

1.8°F/sec could be achieved for the ha If-time-temperature cooling

rate.

One of the interesting conclusions reached from examination

of the possible quenchants was that only slight variations in

cooling rate could be obtained by specific modifications to the

quench medium. Substantial changes could only be produced by

replacing the medium itself. For example, by employing heated

rather than room temperature oil as the coolant, only a 1.74°F/sec

28

change in cooling rate could be obtained. By increasing the flow

rate of air over the plate surfaces from normal circulation a

maximum cooling rate increase of only 0.88°F/sec was possible.

This may be compared to the 8.22°F/sec increase achievable by

substituting an oil quench for air cooling.

The further heat treatment required to obtain the base (un-

embrittled) condition of the HY-100 was established by examining

the tempering characteristics of the mechanical properties.

This dependence is summarized in Figure 5 for hardness and Figure

6 for yield and tensile strength. In Figure 7, a widely used

tempering parameter is used to summarize the hardness data. In

Figure 8 the correlation between the various mechanical properties

produced by these treatments is illustrated. With the aid of

this information, it was concluded that the desired yield strengths

of 110 and 135 ksi could be obtained by tempering the quenched

material for 6 1/2 hours at 1200°F and 1100°F respectively.

Tensile tests performed on material receiving this treatment

demonstrated that, indeed, this temper was satisfactory. The

tensile properties of the unembrittled steel are presented in

Table 3, and the toughness properties of the steel are shown in

Table 4.

While examining these data, it should be noted that this

material, while it has a low as-tempered mid-range transition

temperature, also has very low ambient shelf toughness. The

1200°F tempered material has a shelf toughness of 48 ft-lb, while

29

the 135 ksi yield strength material has a shelf toughness of only

33 ft-lb. These values are low for materials in normal pressure

vessel service. On the basis of typical fracture toughness

characterizations, this material would not be expected to tolerate

cracks larger than about 6 in. long and 1.5 in. deep (i.e., not

through wall cracks) without high potential for fast fracture.

Thus 4 inch thick pressure vessels of this steel operated at

one-third of the tensile strength would not meet leak-before-

break fracture criteria.

TEMPER EMBRITTLEMENT

A summary of the information extracted from the impact

testing of the unembrittled and step cooled HY-100 is furnished

in Table 4. The specific impact curves involved, computer

generated from the energy absorbed and lateral expansion data,

are displayed in Appendix A, Figures A1-A4. From an examination

of these results, two important conclusions may be drawn. It is

apparent that the higher strength material, though exhibiting a

higher transition temperature in the unembrittled condition, is

less susceptible to temper embrittlenient as measured by the shift

in the ductile to brittle transition temperature. This finding

has been reinforced by investigations of other low alloy steels

A <- T U- U (60) performed at Lehigh.

Secondly, during the step cool embrittlement that is required

to obtain an immediate evaluation of the susceptibility of the

30

material, a discernible drop in the upper shelf toughness as

measured by both the energy absorbed and lateral expansion is

evident. The deterioration of high temperature impact resistance,

was also perceivable from the evaluation of subsize Charpy tests.

The influence of the kinetics of temper embrittlement on

impact properties is depicted in Figures 9 and 10 and in Table 5.

The accompanying impact curves are reproduced in Appendix A,

Figures A5 through A20. It is apparent from this information

that, at 750°F, the degeneration of impact resistance, character-

istic of temper embrittlement, is slow to develop. Only a slight

embrittlement, a 30°F shift or less, is produced after 1000 hours

at 750°F at both strength levels. In addition, consistent with

the step cooling results, isothermal aging demonstrates that the

135 ksi yield strength material is less susceptible to temper

embrittlement than the 110 ksi yield strength material.

Figures 9 and 10 illustrate that not only is embrittlement a

slow process at 750°F, it is preceded by a modest decrease in the

transition temperature. This characteristic, also recognized in

earlier studies, suggests that while embrittlement is initially

sluggish, other processes resulting in slight toughening will

occur. However, at the higher embrittling rates which accompany

exposure to higher temperatures, on the order of 900°F, or after

prolonged subjection to lower temperatures, greater than 1000 hours

at 750°F, this toughening is overwhelmed by temper embrittlement.

Comparison of the step cool and isothermally embrittled

31

impact data demonstrates several important characteristics of

these particular aging treatments. Although step cooling permits

an immediate evaluation of the severity of embrittlement in a

specific steel, it does not provide information concerning the

rate of embrittlement at a particular temperature, in this case

750°F. What it does provide is an identification of the relative

maximum degree of embrittlement likely to be encountered in a

material, regardless of the service temperature or exposure time.

Consequently, isothermal data, in conjunction with step cooling

information, illustrates that at the higher strength level, not

only is the susceptibility (maximum embrittlement) of the steel

reduced, the rate at which that maximum is reached is also reduced.

While step cooling accelerates the rate and degree of embrit-

tlement, it also influences the upper shelf of the Charpy impact

curve. Although the step cooling tests would indicate a decrease

in the toughness of HY-100 in the ductile temperature regime,

isothermally treated material exhibits no distinct trend (see

Figures 11A and B). Hence, utilization of a step cooling treat-

ment to assess susceptibility may also produce an artificially

poor value of the upper shelf. The cause for the loss in shelf

toughness on step cooling is not entirely clear. It is known that

this steel can undergo several types of aging phenomena in the

900 to 1000°F range included in the step cool thermal cycle but

not encountered in 750°F aging. Thus step cooling is apparently

not a good predictor of 750°F embrittlement phenomena for the

steel. 32

De-embrittlement at 1050°F for 1 hour of the embrittled

steel demonstrated the reversibility of temper embrittlement.

As Table 6 indicates, not only is the extent of embrittlement

less at a strength level of 135 ksi, the extent of the recovery

of impact resistance is greater for a given time at temperature.

Nevertheless, de-embrittlement was considerable for both the

110 and 135 ksi yield strength material.

Confirmation of temper embritt lement by the evaluation of

tensile properties was also attempted. It was determined, how-

ever, that the properties of unnotched specimens tested at room

temperature were unaffected by embrittlement (see Table 7).

Hence, the conventional tension test cannot be utilized to

assess susceptibility. Similarly, it was found that the mech-

anical behavior of notched specimens subjected to -58°F was not

affected (see Table 6). What was observed was a distinct change

in the appearance of the fracture as embrittlement became more

pronounced. Scanning electron microscopy provided corroboration

of this change.

CREEP EMBRITTLEMENT

Before an evaluation of the influence of an applied load on

the embrittlement at 750°F could be performed, it was necessary

to reproduce the impact data of the unembrittled and step cooled

material. With these data, it was not only possible to compare

the impact toughness of the strained HY-100 directly to the base

condition by utilizing subsize specimens, but the effect of speci-

33

men size on the apparent shift in transition temperature could

be determined.

For subsize specimens, the small range of energy

absorbed on impact (0 - 10 ft-lb) did not facilitate an accurate

determination of the mid-range transition temperature. Conse-

quently, to evaluate the impact data accumulated by testing

smaller Charpy specimens, only lateral expansion data were

employed.

The mid-range lateral expansion transition temperature

changes that are associated with embrittlement are presented

in Table 8. Corresponding impact curves are displayed in

Appendix B, Figures Bl through B6. It is evident from this

information, that while the strength dependence of embrittlement

is again emphasized, the extent of temper embrittlement was

found to be substantially less than that obtained from the

standard size specimens. In particular, less than one half

the degree of embrittlement seen in the larger specimens was

evidenced.

Creep embrittlement, characterized by a decrease in the

upper shelf and the associated increase of the brittle to

ductile transition temperature, was not observed. While creep

strain may induce embrittlement, it will also defer the initi-

ation of temper embrittlement. This retardation is particularly

evident in the HY-100 treated to 110 ksi yield strength in the

unembrittled condition. Rather than a delay of embrittlement,

34

there was an actual improvement of the material as characterized

by a substantial decrease in the transition temperature.

OPTICAL AND ELECTRON MICROSCOPY

Metallography

To reinforce the interpretation of the mechanical

tests, optical metallography was performed. A typical optical

micrograph of quenched and tempered HY-100 is presented in Figure

12A. Its scanning electron counterpart is shown in Figure 12B.

This tempered bainitic structure is not only representative of

unembrittled HY-100, but is also illustrative of the temper

embrittled and creep strained steel. Grain boundary void

formation at prior austenite grain boundaries, indicative of

creep embrittled steels, was not observed. Thus, the conclusion

that creep-embr-ittlement, characterized by a reduction of the

upper shelf and increase in transition temperature, did not

occur was supported metallographically.

Electron Fractography

Figure 13 displays a fracture surface common to HY-100

in the brittle condition. Intergranular fracture and secondary

cracking are both evident. The effect of the de-embrittling

treatment is illustrated in Figure 14. As a result of heat treat-

ment at 1050°F for 1 hour, the faceted fracture appearance is re-

placed by microvoid formation and a substantial increase in impact

resistance.

35

In an attempt to understand the upper shelf drop that

was characteristic of the step cooling treatment, several frac-

ture surfaces of Charpy specimens broken in the ductile tempera-

ture region were examined in the unembrittled, step cooled, and

isothermally aged conditions. However, by observation with the

scanning electron microscope, no dissimilarities in fracture

appearance, typically dimpled, were discerned.

Fractographs of temper embrittled notched and unnotched

tensile specimens are presented in Figure 15. While the unnotched

specimens possessed a dimpled rupture surface after significant

plastic deformation, the notched specimens tested at a low

temperature showed evidence of secondary grain boundary fracture

superimposed upon the dimples. At the strain rates employed in

these tests, however, no significant difference in fracture

strength was observed between these two samples.

36

CONCLUSIONS

1. By step cooling, severe embritt lenient, on the order of a

300°F shift in the mid-range transition temperature, was

produced in the HY-100 quenched and tempered to 110 ksi

yield strength. This shift was determined to be approxi-

mately 150°F in the 135 ksi yield strength material.

2. After isothermal aging at 750°F for 5000 hours, shifts

in the transition temperature of approximately 165°F

and 100°F were observed in the 110 and 135 ksi yield

strength material, respectively.

3. The superposition of creep stress equal to one-half of the

quenched and tempered yield strength upon isothermal aging

at 750°F did not induce creep embrittlement after 1000 hrs.

It did delay the initiation of temper embrittlement.

4. HY-100 quenched and tempered to a higher yield strength,

while possessing a larger unembrittled transition

temperature, was less susceptible to temper embrittle-

ment as characterized by the shift in transition temperature,

5. While step cooling provides an assessment of the maximum

embrittlement likely to be encountered in a steel, regard-

less of temperature, it does not provide a determination

of the degree or rate of embrittlement that will develop

at the service temperature.

6. By step cooling, a perceivable drop in the upper shelf

37

was observed. This decrease of high temperature tough-

ness was not observed during isothermal embrittlement.

Examination of Charpy fracture surface with the scanning

electron microscope could not distinguish this behavior.

7. The embrittlement that evolves during isothermal aging

at 750°F was slow to develop and was preceded by a slight

decrease in transition temperature.

8. Step cooled and isothermally aged HY-100 could be

significantly de-embrittled by heat treatment at 1050°F

for 1 hour.

9. By de-embrittlement of step cooled and isothermally

embrittled Charpy specimens, the predominantly inter-

granular fracture at low temperatures was replaced by

microvoid coalescence and a substantial increase in

toughness . '""

10. Optical and scanning electron microscopy revealed that

the structure produced by heat treatment was bainite.

11. Metallography substantiated that creep embrittlement,

characterized by intergranular void formation, did not

occur .

38

Table 1

CHEMICAL ANALYSIS OF HY-100

Heat No. Material C Mn P _S S_i Ni Cr Mo V Co Ti

15164 HY-100 .18 .35 .010 .025 .29 3.16 1.68 .53 .023 .21 .005

supplied by Earle M. Jorgensen Co., Seattle, Washington

Table 2

EFFECT OF QUENCH MEDIA ON CENTER HALF-

TIME -TEMPERATURE COOLING RATE

ave. cooling rate Medium no. of tests (°F/sec.)

Still air 2 0.96 -> ■

Forced air 4 1.65

High capacity forced air 2 1.84

Still oil 3 9.18

Heated oil 2 10.92

Table 3

MECHANICAL PROPERTIES OF HY-100 TEMPERED AT 1100 AND 1200°F

Temperature(°F) Time(hr) Y.S.(ksi) T.S.(ksi) % Elong. % RA

1100 6 1/2 129.7 152.3 17.2 34.7

1200 6 1/2 108.0 123.1 18.1 52.7

39

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41

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91 73 86

Table 6

DE-EMBRITTLEMENT OF HY100 AT 1050°F FOR 1 HOUR

Heat Treatment Test Temperature(°F) % Embrittlement % Recovery ; E.A. L.E. E.A. L.E.

110 ksi yield strength

Step cooled -25 94

1000 hours -50 29

2000 hours -50 44

5000 hours -40 89

135 ksi yield strength

Step cooled -25 69

1000 hours -25 27

2000 hours -25 5 0

5000 hours -4 46

» „ . .^, _ ,, embrittled . ,nn / _ _ % Embritt lement = (1 r—r—r7—7) x 100 (at testing v unembrittledy v 6 .

temperature)

„ „ ,de-embrittled-embrittled . ,nn , /o Recovery = ( .-,.-.,.... _____ \ x IQQ /at testing J unembrittled - embrittled7 v &v temperature)

E.A. = energy absorbed (ft-lbs)

L.E. = lateral expansion (mils)

67 83 69

8 100 100+

42 100+ 100+

40 100+ 100+

42

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43

Tab Le 8

EFFECT OF APPLIED STRESS ON EMBRITTLEMENT

Shift in lateral expansion Condition trans. temp(°F)

110 ksi yield strength

Step Cooled 98

Aged 1000 hr.-no stress 20

Aged 1000 hr.-55 ksi -80

135 ksi yield strength

Step Cooled 76

Aged 1000 hr.-no stress 20

Aged 1000 hr.-67.5 ksi -80

44

TT SIDE

■2.5" 10"- —25"-H 2.5" ■25"

T 1.5" TOP

thermocouple , protection tube/f ASSEMBLED

Figure 1: Cooling Study - Plate Specifications

45

CREEP SPECIMEN

v :N

FABRICATION OF CHARPY IMPACT SPECIMEN BY FRICTION WELDING

01020 ) OHY100) 01020 )

hiH hi'H h1'H EXTENSION CREEP

SECTION EXTENSION

a

a

r=3 AS WELDED

=F O ORIENTATION OF IMPACT SPECIMEN

TO BE REMOVED

Figure 2: Preparation of Impact Specimens from Creep Specimen

46

1-7/8"

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I I I * I I I I I I

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47

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Figure 4: Effect of Plate Thickness on Center Cooling Rate

48

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TENSILE STRENGTH (ksi)

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YIELD STRENGTH (ksi)

260 270 280 290 300 310 320 330 340 350 360 370 Bhn

Figure 8: Correlation of Mechanical Properties Produced on Tempering

52

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■-)

Figure 12A : Optical Micrograph of Quenched and Tempered HY-100. Tempered bainite structure revealed by 2% Nital etch. 500X.

o

s Figure 12B: SEM Micrograph of Quenched and Tempered HY-100 . Tempered bainite structure revealed by 2% Nital etch.

1700X.

59

Figure 13: SEM Fractograph - Step Cool Temper Embrittled. Note that the intergranular fracture and secondary cracking of the Charpy specimen are predominant. 300X.

Figure 14: SEM Fractograph - Step Cool Temper Embrittled Subsequently De-embrittled. Note the ductile fracture surface of the Charpy specimen. "300X.

60

(^

Figure 15: SEM Fractographs of Temper Embrittled HY-100. Dimpled fracture surface of the unnotched tensile specimen (a) is replaced by dimples with evidence of secondary grain boundary cracking in the notched specimen (b). Both at 400X.

61

REFERENCES

1. Brearley, H., Proc. Inst. Auto. Eng., vol. 11, 1916-17 p. 347.

2. Hollowman, J. H., Trans. ASM, vol. 36, 1946, p. 473.

3. Thum, E. E., "Recent Accidents with Large Forgings," Metal Progress, vol. 69, 1956, p. 49.

4. Schaefer, 0., 'Work of the Task Group on Brittle Failure of Large Forgings," ASME paper #55-A-209.

5. Kalderon, D., Proc. Inst. Mech. Eng., vol. 180, 1972, p. 341.

6. Emmer, L. G., Clauser, C. D., and Low, J. R., "Critical Literature Review of Embrittlement in 2 1/4 Cr-1 Mo Steel," Paper presented to PVRC, 1972.

7. Woodfine, B. C, JISI, vol. 173, 1953, p. 229.

8. Woodfine, B. C, JISI, vol. 173, 1953, p. 240.

9. Jolivet, H., and Vidal, G., Rev. Met., vol. 41, 1944, p. 387.

10. Low, J. R., Fracture of Eng. Materials, ASM, 1964, p. 127.

11. Carr, F. L., Nunes, J., and Larson, F. R., ASTM STP 407, 1968, p. 203.

12. McGregor, G., "Fatigue and Fractured Metal," 1959, New York: The Technology Press of the MIT and Wiley, p. 229.

13. Chapman, R. D., and Jominy, W. E., Trans. ASM, vol. 45, 1953, p. 710.

14. Weiss, B. Z., Niedzwiedz, S., and Bauer, M., JISI, vol. 186, 1966, p. 152.

15. Yoshino, K., and McMahon, C. J., Met. Trans., vol. 5, 1974, p. 363.

» 16. Jaffe, L. D., and Buffum, D. C, Trans. ASM, vol. 42, 1950,

p. 604.

17. Carr, F. L., Goldman, M., Jaffe, L. D., and Buffum, D. C, Trans. ALME, vol. 197, 1953, p. 998.

62

18. Cohen, J. B., Hurlich, A., and Jacobson, M., Trans. ASM, vol. 39, 1947, p. 109.

19. Guttmann, M., Krahe, P. R., Abel, F., Amsel, G., Bruneaux, M., and Cohen, C, Met. Trans., vol.. 5, 1974, p. 167.

20. Powers, A. E., JISI, vol. 186, 1957, p. 323.

21. Vidal, G., Rev. Met., vol. 42, 1945, p. 149.

22. Preece, A., and Carter, R. D., JISI, vol. 173^ 1953, p. 387.

23. McMahon, C. J., ASTM STP 407, 1968, p. 127.

24. Ohtani, H., Feng, H. C, and McMahon, C. J., Met. Trans., vol. 5, 1974, p. 516.

25. Low, J. R., Stein, D. F., Turkalo, A. M., and Laforce, R. P., Trans. AlME,-vol. 242, 1968, p. 14.

26. Gould, G. C, ASTM STP 407, 1968, p. 59.

27. Stevens, W., and Balajiva, K., JISI, vol. 193, 1959, p. 141.

28. Shultz, B. J., and McMahon, C. J., ASTM-STP 499, 1972, p. 104.

29. Balajiva, K., Cook, R. N., and Worn, D. K., Nature, vol. 178, 1956, p. 433.

30. Austin, G. W., Entwisle, A. R., and Smith, G. C, JISI, vol. 173, 1953, p. 376.

31. Viswanathan, R., Met. Trans., vol. 2, L971, p. 809.

32. Restaino, P. R., and McMahon, C. J., Trans. ASM, vol. 60, 1967, p. 699.

33. Marcus, H., and Palmberg, P., Trans. AIME, vol. 245, 1969, p*. 1664. \

34. Stein, D., Joshi, A., and Laforce, R. P., Trans. ASM, vol. 62, 1969, p. 776.

35. Inman, M. C, and Tipler, H. R., Acta Met., vol. 6, 1958, p. 73.

36. McLean, D., and Northcott, L., JISI, Vol. 168, 1948, p. 169. <*-

63

37. Smith, C. L., and Low, J, R., Met. Trans./ vol. 5, 1974, p. 279.

38. Guttmann, M., and Krahe, P. R., Mem. Rev. Met., vol. 70, 1973, no. 9.

39. Thomas, W. R., and Chalmers, B., Acta Met., vol. 3, 1955, p. 17.

40. Palmberg, P., and Marcus, H., Trans. ASM, vol. 62, 1969, p. 1016.

41. Cohen, J. B., Hurlich, A., and Jacobson, M., Trans. ASM, vol. 39, £947, p. 109.

42. Schultz, B. J., and McMahon, C. J., Met. Trans., vol. 4, 1973, p. 2485.

43., Wada, T., and Doane, D. V., Met. Trans., vol. 5, 1974, p. 231.

44. Capus, J. M., JISI, vol. 200, 1962, p. 922.

45. Jaffe, L. D., Carr, F. L., and Buffum, D. C, Trans. AIME, vol. 197, 1953, p. 1147.

46. Hurlich, A., Trans. ASM, vol. 40, 1948, p. 805.

47. Keh, A. S., and Porr, W. C, Trans. ASM, vol. 52, 1960, p. 81.

48. Steeb, E. F., and Rosenthal, P. C, Trans. AIME, vol. 212, 1958, p. 706.

4,9. Irani, J. J., May, M. J., and Elliott, D., ASTM STP 407, 1968, p. 168.

50. Aust, K. T., Hannemann, R. E., Niessen, P., and Westbrook, J. H., Acta Met., vol. 16, 1968, p. 291.

.51. Andrew, J. H., and Dickie, H. A., -JISI, vol. 114, 1926, p. 359.

52. Re Hick, J. R., and McMahon, C. J., Met. Trans., vol. 5, 1974, p. 2439.

53. Joshi, A., and Stein, D., Met. Trans., vol. 1, 1970, p. 2543,

54. Joshi, A., and Stein; D., ASTM STP 499, 1971, p. 59.

64

55. Marcus, H., Hacket, W., and Palmberg, P. W., ASTM STP 499, 19.71, p. 90. .,

56. Pellini, W. E., and Quereon, B. R., Trans. ASM, vol. 39, 1947, p. 139.

57. Sachs, G., and Brown, W., ASTM STP 128, p. 6.

58. Matters, R., and Dickenson, C, Trans. ASME, vol. 130, 1958.

59. Buchmann, W., MITT I.V.G.B. No. 65, 1937, p. 358.

60. Clauser, C. D., Emmer, L. G., Pense, A. W., and Stout, R. D., "A Phenomenologocal Study of the Susceptibility to Temper Embrittlement of 2.225% Cr-1% Mo", Paper presented at the 37th midyear meeting of the American Petroleum Institute's Division of Refining, May 11, 1972.

61. Clauser, C. D., Pense, A. W., and Stout, R. D., M.S. Thesis Lehigh University, 1971.

62. Ferrell, D. E., and Pense, A. W., "Creep Embrittlement of 2 1/4% Cr-1% Mo Steel," Paper presented to the PVRC, 1973.

I -

63. Swift, R., Lukens Steel Co., RDR 69-23, Dec. 1969.

64. Woodford, D., and Goldholf, R., Materials Science and Engineering, vol. 5, 1969, p. 303.

65. Kerr, J., Materials Laboratory, Richmond, 1966.

66. Roper, C, Lukens Steel Co., RDR 69-2, Jan. 1969.

67. Swift, R., Lukens Steel Co., RDR 69-22, Dec. 1969.

68. Cottrell, A., LSI Special Report 70, 1961, p. 1.

69. Glen, J., Welder, vol. 32, No. 154, April-June 1968.

70. Roper, C, Lukens Steel Co., RDR 68-12, June 1968.

71. Garafalo, F., FundamentaIs of Creep and Creep Rupture in Metals, Ne,w York: The Macmillan Co., 1965.

72. Hale, K., f'An Electron Microscope Study of Changes in a 2 1/4% Cr-i% Mo Superheated Tube Steel During'Tempering and Creep," 4th International Conference on Electron'Micro- scopy, p. 650.

65

73. Bruscato, R., ASME Publication No. 71, Presented at the Petroleum Mechanical Engineering—Underwater Technology Conference, Sept. 19-23, 1971.

66

APPENDIX A: CHARPY CURVES - TEMPER EMBRITTLED HY-100

CD

I

60

50

— 40

03 o CO m tx >- 0C

LU

30

20

10

•*y - "-"T

HY100 - STD. CHflRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

UNEMBRITTLED

—i r ■ — T — T"1

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

CO

CO

cr X LU

cr LU

•»u HY100 - STD. CHflRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

UNEMBRITTLED

I "■ -

x / X /x

/ x

—r r r— — r——

x

—1 1 1

X

30

/"^ X X

-

o

20

x/

/ *

'x -

10 Z * X /

-

n X 1 1 1

-350-300-250-200-150-100-50 0 50 100 150 200 250 300 350

,.. TEST TEMPERATURE (DEG F)

Figure Al: L10 ksi yield strength - Unembr itt led

A-1

CD _J

I I—

o LU CD

o CO en cr

LU

40

30 -

20 -

10 -

HY100 - STD. CHflRPY — i—

' -1 1 —1— 1 r 1

QUENCHED

110 KSI

♦ TEMPERED TO YIELD ST. X X

STEP COOLED X .

X X

- *.

-

-

x ___,^^

X /

~ X

1

X X

-200-150-100 -50 0 50 100 150 200 250 300 350 400 450 500

TEST TEMPERATURE (DEG F.)

CO

CO

cr. X LU

cr ex. LU

cr

40 cCs HY100 - STO. CHflRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

30 STEP COOLED

X X /

/ *

20 x/

/x

X

e- "

10

X __--'

X

•

' X

1 1 —M i- n X X

-200-150-100 -50 0 50 100 150 200 250 300 350 400 450 500

TEST TEMPERATURE (DEG F)

Figure A-2; 110 ks i yield strength - Step Cooled

A-2

CD _l

I

a LU CD an o CO CD cr >- C3 QC LU

40

30

20

10

CO

o CO z ex a. X LU

cr ac

HY100 - 5TD- CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

UNEMBRITTLEO

x -JC-

-350-300-250-200-150-100-50 0 50 100 150 200 250 300 350 400

TEST TEMPERATURE ( DEG F)

30

20

10

HY100 - 5T0. CHflRPY QUENCHED ♦ TEMPERED

TO 135 KSI YIELD ST.

UNEM8RITTLE0

-350-300-250-200-150-100-50 0 50 100 150 200 250 300 350 400

TEST TEMPERATURE (DEG F)

Figure A-3: 135 ksi yield strength - Unembrittled

A-3

30

CD

I

— 20 - o UJ in

o LO CD CX

>- CD QC UJ

UJ

10

0

1 r 1 1 1 —-t —

HY100 - STD. CHARPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

X

X

STEP COOLED

x/ X

At.

-

— /y. -

X

X"/ X ^^-^

X

350-300-250-200-150-100-50 0 50 100 150 200 250 300 350 400 450

TEST TEMPERATURE (DEG F)

30

tS)

CO

CL Q_ X UJ

CX ce UJ

20

10 -

0

HY100 - STD. CHRRPY » — 1— , t — 1 1—

QUENCHED + TEMPERED TO

135 KSI YIELD ST.

STEP COOLED X —

X .-—' ~~ x

- X

X j

s x

X X

-

x . yS X

__2——""'^^ ' ' ■ 1—H-I ■ 1 . . 1

-350-300-250-200-150-100-50 0 50 100 150 200 250 300 350 400 450

TEST TEMPERATURE (DEG F)

Figure A-4: 135 ks i yield strength - Step Cooled

A-4

50

40 CD _J

I I— U_

P 30 CO cc CD co CO cr >- CD ct: UJ

20

10

0.

HY100 - STD. CHARPY QUENCHED + TEMPERED

TO X X

- 110 KSI YIELD ST. X X

AGED 1 HOUR */ X

AT 750 F Y X

/ * ,

x/x -

X /

X /

- ^^ -

X

J 1 1

'

350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

30

co

-z. CD

CO

(X a_ x UJ

ex

UJ (— a;

20

10

HY100 - 5TD- CHflRPY ?— — 1— — 1 1 - r — — I 1

QUENCHED

., 110 KSI AGED

AT

♦ TEMPERED TO x

YIELD ST. / 1 HOUR */ 750 F /

' X

X

X

-

/ *

/x x / x/ / X

X /

X

X

-350-300-250-200-150-100-50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

Figure A-5: 110 ksi yield strength - Aged 1 hour

A-5

CD _J

I

□ LU CO ce o CD cr >~ Od LU

LU

50

40

30

20

10

HYIOO - 5T0. CHflRPY QUENCHED + TEMPERED

TO 11.0 KSI YIELD ST. AGED 10 HOURS

AT 750 F

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

to

en 2 d Q_ X LU

cr ce LU

40

30

20

10

HY100 - 5TD- CHflRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

AGED 10 HOURS AT 750 F

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

Figure A-6: 110 ks i yield strength - Aged 10 hours

A-6

CD

I

o LU CO QC a m CO ex >- o QC LU

LU

a. Q_ X LU

(X a; LU

60

50

- 40

30

20

10

HY100 - 5TD. CHARPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

AGED 30 HOURS PT 750 F

C

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

50

in d 40

^ 30 en

20

10

HY100 - STQ. CHARPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST. AGED 30 HOURS

AT 750 F

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

Figure A-7; 110 ksi yield strength - Aged 30 hours

A-7

CD

I

o UJ m oc o to 03 cr

>- Q: UJ

SO

40 -

30 -

20 -

10

Q

HY100 - 5TP- CHARPY QUENCHED + TEMPERED

TO 110KSI YIELD ST. AGED 100 HOURS

AT 750 F

350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

CO

o »—I

CO

cr Q_ X UJ

cr a: UJ

N50

40

£ 30

HY100 - STQ. CHARPY QUENCHED ♦ TEMPERED

TO - 110 KSI YIELD 5T.

AGED 100 HOURS AT 750 F

20

10

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE ( DEG F)

Figure A-8: 110 ksi yield strength - Aged 100 hours

A-8

CD

CD OC CD CO 00 cr

50

40

30

20

LU 10 z LU

HY100 - 5T0. CHRRPY QUENCHED ♦ TEMPERED

TO 110 KSI YIELD ST. AGED 300 HOURS

AT 750 F

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

40

CO

O »—^ CO 2 d Q_ X

cr Of LU

10 -

HY100 - 5T0. CHRRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

30 h AGED 300 HOURS AT 750 F

20 -

-350-300-250-200-150-100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (DEG F)

' Figure A-9: 110 ksi yield strength - Aged 300 hours

A-9

50

-J 40 - CD

7 i—

o LU CD

o co GO cr. >- o LU

30 -

20

10

HY100 - 5T0. CHARPY QUENCHED ♦ TEMPERED

TO 110 KSI YIELD ST. AGED 1000 HOURS

AT 750 F

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

50

co d 40

CO

<x Q_ X LU

cr ce LU \— ex

30

20 -

10

HY100 - STD. CHARPY 1— 1 f

QUENCHED + TEMPERED TO

- 110 KSI YIELD ST. AGED 1000 HOURS

AT 750 F

X *

/ X

X X

X X

X

-

- / *

-

x /> .X, ^/ -

X ■ 111 .1 1 1— 1 I 1 ' ■

-350 -300 -250 -200 -150.-100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEC F)

Figure A-10: LIO ks i yield strength - Aged lOOO hours

A-10

CX

CD

I

a LU CD Ql CD m CD a. >- o or- UJ

LU

60

50 -

40

30

20

10

HY100 - 5TP. CHARPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST. AGED 2000 HOURS

FIT 750 F

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

50

o ^ 30 CD Z ex. Q_ X LU

d

LU

20

10

HY100 - STD. CHRRPY ^^ QUENCHED + TEMPERED CO TO _J 40 - 110 K5I YIELD ST. x: AGED 2000 HOURS *-^ AT 750 F

-350 -300 -250 -200 -150-100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

Figure A-11: 110 ks i yield strength - Aged 2000 hours

A-ll .

CD

I

CD ai CD CO 03 CX

>- CD oc LU

50

40

30

20

10

HY100 - 5T0. CHARPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST. AGED 5000 HOURS

AT 750 F

-150 -100 -50 0 50 100 150 200

TEST TEMPERATURE (DEG F) 250 300 350

50

CO

d 40

CO

cc Q_ X UJ

cc en LU \—

30

20 -

10 -

HY100 - STD. CHARPY T ■■ "I 1

QUENCHED + TEMPERED TO

- 110 KSI YIELD ST. X * - AGED 5000 HOURS X -—

AT 750 F X

X j.

/ X

& x / / X

/ -

y/ X ,,

x-; , i 1 1 . ' 1 . . -150 -100 -50 0 50 100 150 200

TEST TEMPERATURE (DEG F) 250 300 350

Figure A-12: 110 ksi yield strength - Aged 5000 hours

A-12

<3

CD _l

I I— u_

CD Ctl O CO CD <X

>- CH LU

40

30

20 -

10 -

HY100 - 5TD. CHflRPY r 1 1 1 1 T 1 1

QUENCHED + TEMPERED TO

135 KSI YIELD ST. AGED 1 HOUR

AT 750 F X s^

X

X

yC x

x/

~ /x

—

e y /x X

X ^/

' ■ «

X

1 1 1 I

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

30

CO

o CO

ex Q_ X

ex ct: LU

20 -

10

—, ,_ HY100 - 5TD. CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

AGED 1 HOUR AT 750 F

-350 -300 -250 -200 -150-100-50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

Figure A-13: 135 ksi yield strength - Aged 1 hour

A-13

CD _J

I I— U_

CD QZ O CO CO CX

>-

QC UJ zz. UJ

40

30 -

20

10

0

HY100 - STD. CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST. AGED 10 HOURS

AT 750 F

i -1

V

1 r

X X

1

X T 1

/ X X

- /x -

x y x x x ^s"

X

350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEC F)

30

CO

O »—* CO

CX Q_ X UJ

CX

UJ

20 -

10 -

HY100 - STD- CHflRPY —i r- "■ "1 T .... r

QUENCHED + TEMPERED TO

135 KSI YIELD ST. AGED 10 HOURS

AT 750 F

X

x/

X

X X

X /

/x

- / / X

-

x ^^"x

1 1 1

X

"350 ^300 -250 -200 -150 -100 -50 0 50 100

TEST TEMPERATURE (DEG F)

150 200 250

iFigure A-14: 135 ks 1 'yield strength - Aged 10 hours

A-14

en _i

i i— Li-

en Od CD CO CD CX

>- ce LU

LU

40

30

20

10

HY100 - STQ. CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST. RGED 30 HOURS

AT 750 F

-i 1 p r-

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

30

to

CO

CX Q_ X LU

cr

LU

20

10

HY100 - STD. CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

AGED 30 HOURS AT 750 F

-350 -3UU -2bLT^2bo -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

i

Figure A-15: 135 ksi yield strength - Aged 30 hours

A-15

CD _J

I I— U_

CO ac CD in CD ex.

LU

50

40

30

20

10 -

HY100 - 5TD. CHARPY T 1 —i 1

QUENCHED + TEMPERED TO

- 135 KSI YIELD ST. AGED

AT 100 HOURS 750 F

X x x

- X yT'

y x

-

- x .

X j/

/x X

a

-

- x/^ -

X x

1 !-.„ • ' ' — -350 -300 -250 -200 -150-100 -50 0 50 100 150 2p0 250

TEST TEMPERATURE (DEG F)

30

tn

CD

(Si ■z. QZ. Q_ X LU

<X QC LU

20

10

HY100 - 5T0. CHARPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST. AGED 100 HOURS

RT 750 F

X

/x

-T 1 1 1 1 1—

X i

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

Figure A-16: 135 ks i yield strength - Aged 100 hours

A-16

40 HY100 - 5TD- CHflRPY QUENCHED'+ TEMPERED

TO 135 KSI YIELD ST. AGED 300 HOURS

AT 750 F

-350 -300 -250 -200 -150-100 -50 0 50 100

TEST TEMPERATURE (DEG F)

150 200 250

30

en

<x Q_ X LU

a. UJ

20 -

10 -

HY100 - STD. CHflRPY r

QUENCHED + TEMPERED TO

135 KSI YIELD ST. AGED 300 HOURS

AT 750 F

X

/* X

/x y *

- . /

x:/

x/ •4

x ^—^-^^ X

u 1— ■ • • 1 .A J

-350 -300 -250 -200 -150-100-50 0 50 100 150 200 250

TEST TEMPERATURE [DEG F)

Figure A-17: 135 ksi yield strength - Aged 300 hours

A-17

CD _1

I

ao ac CD CO CO cr >- ce LU

40

30

20

10

1 r* 1 1—

HYI00 - 5TD. CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

AGED 100D HOURS AT 750 P

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

30

CO

co cx Q_ X LU

CX CK

20

10 -

0

HY100 - ST0. CHPRPY r ■■ i T— r

QUENCHED + TEMPERED TO

135 KSI YIELD ST. flCED 1000 HOURS

AT 750 F

X X

* ,--»r—

x s'

/x

X /

x it/ /x

— 1 i i

350 -300 -250 -200 -150-100 -50 0 50 100

TEST TEMPERATURE (DEG F) 150 200 250

Figure A-L8: 135 ksi yield strength - Aged 1000 hours

A-18

_J I

I—

00 Q£ O CO CD CL

>- ct: LU

40

30

20 -

10

HY100 - STO. CHflRPY —I 7 r 1 1— 1 —1

■ QUENCHED + TEMPERED TO

135 KSI YIELD ST. - flGEO 2000 HOURS

AT 750 F

X /

X

X X

X

X /x

X y x

X

u

-

X

-350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

to

o

co

cr

X UJ

ex

LU

40

30

20

10

HY100 - STD. CHBRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

h flCEO 2000 HOURS AT 750 F

0 350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEG F)

Figure A-19: 135 ksi yield strength - Aged 2000 hours

A-19

CD _J

I

□ UJ CD ai o co CO en >-

LU

40

30

20 -

10 -

HY100 - STD- CHflRPY t ! ■ 1 1 ■ 1 -

QUENCHED + TEMPERED TO

135 KSI YIELD ST. X

- AGED 5000 HOURS X -

AT 750 F X

^^x

- /x -

X S*

x^x . -

X

I , 1 1 1 1 1

-150 -100 -50 0 50 100 150 200

TEST TEMPERATURE (DEG F)

250 300 350

30

cO

CO

ex Q_ X UJ

cr ct: UJ (— ex

20

10

HY100 - STD. CHflRPY 1 1 ' T- VI '

QUENCHED + TEMPERED TO

135 KSI YIELD ST. AGED 5000 HOURS

HT 750 F X

X

—' X

*/ X X

^

./OC

/ X x /

X-"^ X 1 1 1 1 i —1 1

-150 -100 -50 0 50 100 150 200

TEST TEMPERATURE (DEG F) 250 300 350

Figure A-20: 135 ksi yield strength - Aged 5000 hours

A-20

APPENDIX B: CHARPY CURVES - CREEP EMBRITTLED HY-LOO

20

CD _J

I \—

o LU CD C£ CD CD CO cr >- CD Cn LU

10

0

HY100-SUB5IZE CHflRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

UNEMBRITTLED

350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE fDEG F)

-O

_i

to

(X Q_ X LU

az en LU

40

30 -

20

10

HY100-SUBSIZE CHflRPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

UNEMBRITTLEO ( x X

X X

X X/ X

X /X -

/ x X X. X

X /

/ X X / X -

l&x* X ^sX

^~-^v< ■

X

-350 -300 -250 -200 -150-100-50 0 50 10(0 150 200 250

TEST TEMPERATURE (DEG F) A

Figure B-l: 110 ksi yield strength - Unembrittled

A-21

10

m <r o o> m <

to

HYIOO-SUBSIZE CHARPY QUENCHED TEMPERED

TO 110 KSI YIELD ST.

STEP COOLED

-250 -200 -150 -100 -50 0 50 100 150 200 250 300 350

TEST TEMPERATURE (°F)

-250 -200 -50 0 50 100 150

TEST TEMPERATURE (°F)

\

Figure B-2: 110 ks i yield strength' - Step Cooled

A-22

10

a UJ CD

8 CO CO < >- e> DC Ul z Ul

HYIOO-SUBSIZE CHARPY QUENCHED * TEMPERED

TO 110 KSI YIELD ST.

CREPT 1000 HOURS AT 750F

-350 -300 -250 -200 -150 -100 -50 0 50

TEST TEMPERATURE CF)

100 150 200 250

20

5 2 z o CO z

X

< Id

10

-I 1-

HYIOO-SUBSIZF CHARPY QUENCHED + TEMPERED

TO 110 KSI YIELD ST.

CREPT 1000 HOURS AT 750 F

-350 -300 -250 -200 •150 -100 -50 0 50

TEST TEMPERATURE (°F)

100 150 200 250

Figure B-3: 110 ksi yield strength - Crept 1000 hours

A-23

(

20

CD _J

I I— li-

en UJ CD on CD en CO ex >■

CD C£

10 —

1 1 1— —1

HY100-SUB5IZE CHARPY —i 1 —i—

_ ._ 1 T—

QUENCHED + TEMPERED TO

135 KSI YIELD ST. UNEMBRITTLED

X

x >r

x'x

X

X X

—" X X X

X

x x__^— i . i— ■

. . . -350 -300 -250 -200 -150-100 -50 0 50 100 150 200 250

TEST TEMPERATURE I DEC F)

m

to cr X UJ

cr

40

30 -

20

10 -

HY100-SUBSIZE CHflRPY 1^— —r ■ r ! 1 '— I - —r '

QUENCHED + TEMPERED TO

135 KSI YIELD ST. UNEMBRITTLED X

X

X

X X

X K

X

-

X

X/ic X /x

7 X

-

" X

X -

_^^X

0 -350 -300 -250 -200 -150 -100 -50 0 50 100 150 200 250

TEST TEMPERATURE (DEf> P-)

Figure B-4: 135 ksi yield strength - Unembrittled

A-24

20

CD _l

I I—

o m a: o CO CO GC

>- o UJ

10

HY100-5UB5IZE CHRRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

STEP COOLED

-ZOO -150 -100 -50 0 50 100 150

TEST TEMPERATURE iDEG F) 200 250

30

cO

O

CO -z. CX Q_ X

CX ct: LU

20 -

10

Q

MY100-SUBSIZE CHRRPY 1 1 1 1 '

QUENCHED + TEMPERED TO

135 KSI YIELD ST.

STEP COOLED x/

X X

X

X

X / X / x/ / X

X ,

— X X - X

X

X X

^^tr^^ x

1 1 I I

X MX X X X

1 1 1

200 -150 -100 -50 0 50 100 150

TEST TEMPERATURE (DEG F) 200 250

^1

Figure B-5: 135 ksi yield strength - Step Cooled

A-25

20

CD —I

I

CD QC O CO CD cr >-

cc

10

HY100-SUBSIZE CHflRPY f "- —1 1 '

QUENCHED + TEMPERED to

135 KSI YIELO ST. CREPT 1000 HOURS

AT 750 F •'

— X X

X —

■/ x

*/

x^^/x

* i I I I 1 1 1 I

-200 -150 -100 -50 0 50 100

TEST TEMPERATURE (DEG F)

150 200 250

cD

CO

a. o_ X LU

(X ce

40

30

20

10 -

HY100-SUBSIZE CHflRPY QUENCHED + TEMPERED

TO 135 KSI YIELD ST.

- CREPT 1000 H0UR5 AT 750 F I

■ i - i i i -- r-

/ X

x/

x /

r i , i ' ■ i i ■

-200 -150 -100 -50 0 _ 50 100

TEST TEMPERATURE (DEG F)

150 200 250

Figure B-6: 135 ksi yield strength - Crept 1000 hours

A-26

VITA

Courtney Jay Hill, son of Raymond J. and Jean Hill, was born

in Baltimore, Maryland, on August 5, 1952. A graduate of McDonogh

School, Mr. Hill entered Lehigh University in September, 1970.

Upon graduation in May, 1974, he was awarded the degree of Bachelor

of Science in Metallurgy and Materials Science with highest aca-

demic and interdepartmental honors. In the summer of 1974, he

was admitted to the graduate school of Lehigh University where he

served as a research and teaching assistant in the Department of

Metallurgy and Materials Science.

Mr. Hill is a member of Tau Beta Pi, The American Society

for Metals, and The Metallurgical Society of AIME.

67