A 1 1 1

2 1405DS NBS

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„ NBSIR 83-1680

INTERIM PROGRESS REPORT:

FRACTURE TOUGHNESS OF STEEL

WELDMENTS FOR ARCTIC STRUCTURES

T. L. Anderson

H. I. McHenry

National Bureau of Standards

U.S. Department of Commerce

Boulder, Colorado 80303

December 1982

56

3-1680

98a

NBSIR 83-1680

INTERIM PROGRESS REPORT:

FRACTURE TOUGHNESS OF STEEL

WELDMENTS FOR ARCTIC STRUCTURES

Fracture and Deformation Divison

National Measurement Laboratory

National Bureau of Standards

U.S. Department of Commerce

Boulder, Colorado 80303

December 1982

Sponsored by

U.S. Department of Interior

Minerals Management Service

12203 Sunrise Valley Drive

Reston, Virginia 22091

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

T. L. Anderson

H. I. McHenry

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

CONTENTS

Page

1. INTRODUCTION 1

2. TECHNICAL APPROACH 2

3. CURRENT FRACTURE TOUGHNESS REQUIREMENTS 3

3.1 Pressure Vessels and Piping 4

3.2 Ships 5 3.3 Bridges 5

3.4 Summary Comments 6

4. EXPERIMENTAL PROCEDURE 6

4.1 Test Material 6

4.2 Tensile Tests 6

4.3 Charpy V-Notch Impact Tests 7 4.4 Fracture Toughness Tests 7

5. RESULTS AND DISCUSSION 10

5.1 Tensile Tests 10

5.2 Charpy- Impact Data 10

5.3 Fracture Toughness 11

5.4 Predicting the Effect of Constraint on the Ductileto-Brittle Transition 12

5.5 Detecting the Onset of Tearing 13

5.6 Estimation of J from CMOD 14

5.7 Relationships Between J and CTOD 15

5.8 The Eta Factor 16

6. SUMMARY AND CONCLUSIONS 17

REFERENCES 18

APPENDIX 50

i i i

List of Figures Page

1. The ducti le-to-brittle transition of steel showing brittle (linear

elastic), elastic-plastic, and ductile (fully plastic) regions. 31

2. Location of Charpy V-notch specimens for weld qualification tests. 32

3. Single-edge notched bend (SENB) specimen used to obtain fracture 33

toughness values.

4. Orientation of SENB specimens with respect to the rolling direc- 34

tion. Charpy impact specimens were also prepared with this

orientation.

5. Photograph of the test machine and three-point bend fixture. 35

6. Notation for reporting critical CTOD values [13]. 36

7. Tensile properties of ABS grade EH36 steel as a function of 37

temperature.

8. Charpy impact transition curve for ABS grade EH36 steel. 38

9. Critical CTOD for cleavage as a function of temperature and speci- 39 men thickness for ABS grade EH36 steel. W = 25.4 mm (1.0 in).

10. Critical J for cleavage as a function of temperature and specimen 40

thickness for ABS grade EH36 steel. W = 25.4 mm (1.0 in).

11. Critical CTOD for cleavage as a function of temperature and crack 41

length for ABS grade EH36 steel. W = 25.4 mm (1.0 in).

12. Critical J for cleavage as a function of temperature and crack 42

length for ABS grade EH36 steel. W = 25.4 mm (1.0 in).

13. Critical CTOD for the onset of tearing as a function of temperature 43

and specimen geometry for ABS grade EH36 steel. W = 25.4 mm (1.0 in).

14. Critical J for the onset of tearing as a function of temperature and 44

specimen geometry for ABS grade EH36 steel. W = 25.4 mm (1.0 in).

15. Comparison of predicted and experimental CTOD values as a function 45

of temperature and specimen thickness.

16. Computer plots of the first and second derivatives of q with 46

respect to V. A sharp drop in the first derivative generally

indicates tearing.

17. Comparison of J estimates from equations 8 and 9. 47

18. The ratio of the plastic components of J and CTOD and an estimate 48

of the plastic m factor as a function of displacement.

19. The overall eta factor (n ) as a function of a/W and displacement. 49

The elastic eta factor curve [17] is superimposed for comparison.

v
List of Tables Pane

1. Charpy V-Notch Requirements for Steel Plates in OCS Platforms. 20

2. Charpy V-Notch Requirements for Fracture-Cri tical Members in 21

Arctic Bridges.

3. Charpy V-Notch Requirements for Pressure Vessels and Piping. 22

4. Chemical Composition of the ABS Grade EH36 Steel. 23

5. Test Matrix for Fracture Toughness Tests on SENB Specimens. 23

6. Tensile Properties of the ABS Grade EH36 Steel. 24

7. Charpy V-Notch Impact Data for ABS Grade EH36 Steel. 25

8. Summary of Fracture Toughness Tests on Five Geometries of 26

SENB Specimens.

a. B = 0.5 W = 12.7 mm 26

b. B = W = 25.4 mm 27

c. B = 15 W = 38.1mm 28

d. B = W = 25.4 mm 29

e. B = W = 25.4 mm 30

vi
FRACTURE TOUGHNESS OF STEEL WELDMENTS

FOR ARCTIC STRUCTURES

T. L. Anderson

H. I. McHenry

This report summarizes the progress in the development of fracture

criteria for steel weldments in arctic structures. Tensile, Charpy-impact

,

and fracture toughness properties have been measured as a function of temperature for a 25.4 mm (1 in) thick plate of normalized steel. Fracture toughness

tests were performed on five geometries of single-edge notched bend (SENB)

specimens. Critical values of the J-integral and the crack-tip opening

displacement (CTOD) were computed and plotted versus temperature. The ducti 1 e-to-bri ttl e transition temperature increased with increasing specimen

thickness, and crack length. The effect of specimen geometry on fracture

toughness is attributed to changes in crack-tip region constraint with geometry.

Initial attempts to model this behavior have been moderately successful.

Various aspects of the SENB fracture toughness test are being examined.

Preliminary results indicate that the J-integral can be accurately measured

with either load-line displacement or mouth-opening displacement measurements.

It may therefore be possible to measure both J and CTOD with a single clip

gage. The relationship between J and CTOD has been investigated. The ratio

of J to CTOD is a function of yield strength, displacement and work-hardening

rate. The eta factor, which is a dimensionless constant used to relate the

J-integral to energy absorbed by the specimen, was found to be independent of

crack length for SENB specimens with a/W ranging from 0.19 to 0.75.

Key words: arctic structures; crack-tip opening displacement; ductile- tobrittle transition; elastic-plastic fracture; J-integral; mechanical properties; structural steels; toughness.

1. INTRODUCTION

The development of oil and gas resources in the Arctic will require the

construction of fixed offshore platforms capable of operating safely at

temperatures of -40°C and possibly lower. The platforms are likely to be

welded steel structures because of steel's relatively low cost, ease of

- 1

-
fabrication, high strength, and the availability of grades that are highly

fracture resistant at low temperatures. The arctic environment is potentially

hazardous from a structural integrity standpoint because steel weldments have

increased susceptibility to brittle fracture at low temperatures . Thus,

material toughness criteria are needed to ensure that steels with adequate

fracture resistance at low temperatures are used to build these platforms.

The objective of the present study is to develop a quantitative basis for

specifying the fracture toughness of steel weldments for arctic structures.

The present report summarizes the overall program plan and the progress

to date. The technical approach, which is based on elastic-plastic fracture

mechanics, is summarized in the next section. A review of current code

requirements for the fracture toughness of steels used in arctic structures is

given in Section 2. The experimental procedures, the results and discussion,

and the summary and conclusions are given in Sections 4, 5, and 6, respectively. A literature review covering the basic concepts of fracture mechanics

and recent research pertinent to the present study is included as an appendix.

2. TECHNICAL APPROACH

The principles of fracture mechanics [1,2] provide the technical basis

for this investigation. Fracture mechanics is the study of the influence of

loading, crack size, and structural geometry on the fracture resistance of

materials containing cracks. There are two main ideas in fracture mechanics.

First, fracture occurs when the driving force for fracture, a function of

stress and flaw size, exceeds the material's resistance to fracture, referred

to as the fracture toughness. Second, fracture toughness is a geometryindependent material property; i.e., a simple laboratory specimen and a large

structure both fracture at the same critical value of driving force. Both of

these ideas work quite well for materials that behave in a linear elastic

manner; i.e., fracture occurs prior to extensive plastic deformation. Recent

progress in ductile fracture mechanics suggests that these ideas are also

valid for materials that fracture in a fully ductile manner [3].

Contrary to common assumptions, there is not a geometry- independent

fracture criterion for the transition region between linear elastic and

fully-ductile fracture. This transition region, identified as elastic-plastic

in Fig. 1, is the region of interest for steel structures at low temperatures.

The geometry dependence exists because steel undergoes a ductile- to-brittle

- 2 -
transition as temperature is reduced. In the temperature range of the transition, different fracture mechanisms may occur depending on the extent of

plastic flow near the crack tip. As plastic flow decreases, the micromode of

fracture changes from ductile tearing to brittle cleavage. Plastic flow can

be decreased by lowering the temperature or increasing the strain rate, both

of which increase the yield strength. Plastic flow can also be limited by

geometrical constraint, i.e., configurations where the surrounding elasticmaterial limits plastic flow near the crack tip. Constraint increases with

thickness (the well-known size effect) and, in bending, with crack depth-tothickness ratio.

The fracture behavior of steels at low temperature is studied using the

principles of elastic-plastic fracture mechanics (EPFM). Both the driving

force and the resistance to fracture can be expressed in terms of an EPFM

parameter, either the crack tip opening displacement (CTOD) or the J-integral

(J). Fracture occurs when the applied CTOD (or J) exceeds a critical value of

CTOD (or J). The critical value of CTOD, denoted 6 , or of J (J

f

) is the

fracture toughness of the material . Unfortunately , in the temperature range

of the ductile-to-brittle transition, 6 c and J

c are not geometry independent.

Thus, one of the main ideas of fracture mechanics, the applicability of

laboratory fracture toughness data to practical structures, is not valid.

The approach being used to develop fracture criteria for arctic structures is based on EPFM. Fracture toughness of a C-Mn steel plate is measured

as a function of geometrical constraint at temperatures throughout the ductile

to-brittle transition. Constraint is varied by changing the crack-length

specimen-width ratio and the thickness of single edge notch bend (SENB)

specimens. In each test, fracture toughness is measured by the two principal

EPFM parameters, CTOD and J. The effect of constraint on fracture toughness

is modeled in an attempt to predict the temperature shifts in the ductile-tobrittle transition curve with changes in constraint. A suitable method to

account for constraint may eventually be used in conjunction with the Burdekin

Dawes [4,5] design curve to assess the f i tness-for-service of steel weldments.

3. CURRENT FRACTURE TOUGHNESS REQUIREMENTS

In 1979, the United States Geological Survey published the "Requirements

for Verifying the Structural Integrity of OCS Platforms" [6]. These

- 3-
requirements include state-of-the-art performance standards which must be met

in designing, fabricating, and installing OCS platforms, including those on ice or gravel islands in the arctic. The objective of the requirements is to ensure that offshore oil and gas platforms and other structures have a high

probability of surviving the environmental and operational conditions to which

they may be exposed.

To ensure satisfactory fracture resistance, the USGS provided the guidelines shown in Table 1 for specifying minimum levels of Charpy V-notch (CVN)

impact toughness for low and intermediate strength steels. The CVN toughness

guidelines for intermediate strength (310 to 415 MPa yield strength) steel

plates is 34 J in the longitudinal orientation. The test temperature for

steels in the most severe applications is 30°C below the service temperature.

Testing at 30°C below the service temperature is not practical in arctic

structures because economical steels with sufficient toughness at -70°C are simply not available. Thus, the requirements state: "For service temperatures

less than -30°C, test temperatures (for CVN testing) should be specially

considered .

The specification of an energy level and test temperature for CVN testing

varies for different types of structures that are subjected to low temperatures during service. In the following subsections, the fracture criteria

specified for pressure vessels and piping, ships, and bridges are reviewed.

3. 1 Pressure Vessels and Piping

The toughness requirements for pressure vessels (ASME Section VIII, [7])

and piping (ANSI and B31.3 [8]) are a function of strength level, deoxidation

practice, and heat treatment. Minimum CVN values for carbon and low alloy

steels are specified in terms of average absorbed energy at a specified

temperature as shown in Table 2. The toughness requirements for higher

strength steels, steels heat treated to enhance strength, and stainless steels

are specified in terms of notch ductility as measured in a CVN test, 0.38 mm (0.015 inch) lateral expansion is required. The test temperature for all

grades is the design temperature or the minimum temperature at which pressure

will be applied, whichever is lower. A lower test temperature must be used

when subsize tests are conducted and the subsize width is less than 80% of the

material thickness; e.g., the test temperature for half- and quarter-size CVN

specimens are reduced by 11 and 28°C, respecti vely

.

- 4 -
For welded pressure vessels, the CVN toughness of the weld metal and the

heat-affected-zone (HAZ) must equal or exceed the toughness values specified

for the base metal at the same (or lower) test temperature. Tests are con- ducted on the weld procedure qualification test plates and the production test

plates. Weld metal specimens are notched perpendicular to the surface with

one face of the specimen located within 1.5 mm of the surface. The HAZ

specimens are etched before notching and the notch is cut approximately normal

to the surface in such a manner as to include as much HAZ material as possible

in the resulting fracture.

3.2 Shi ps

The class of ship that have service temperatures that approximate arctic

conditions are the liquified natural gas (LNG) carriers. In these ships,

significant portions of the hull are cooled by the cryogenic cargo to temperatures ranging from 0 to -46°C [9]. To avoid brittle fracture, the U.S. Coast

Guard (USCG) [10] and the American Bureau of Shipping (ABS) [11] and comparable regulatory authorities around the world have established CVN requirements

on the steels and welds subjected to low temperatures. The CVN tests are conducted at 5.5°C below the minimum service temperature. The specimens are

cut transverse to the rolling direction and to the weld axis with notches

normal to the plate surface. Each plate is tested and welds are tested from

the weld procedure qualification plates and from the production test plates.

Three specimens are tested in each of the following notch locations: base

metal; centered in the weld metal; on the fusion line; and in the HAZ at 1, 3,

and 5 mm from the fusion line. The impact specimen locations are summarized

in Fig. 2. The average CVN value must equal or exceed 27.5 J (20 ft-lb); the

minimum value for one specimen is 18.6 J (13.3 ft-lb).

3.3 Bridges

The toughness requirements for arctic bridges have been established by

the American Association of State Highway Transportation Officials

(AASHTO) [12]. The requirements are applicable to fracture critical members,

which are defined as those tension components of a bridge whose failure would

be expected to result in the collapse of the bridge. The CVN values specified

are summarized in Table 3. For low and intermediate strength steels, the test

temperature is 39°C above the minimum service temperature, i.e., a bridge

- 5 -
design for -51°C service uses steels tested at -12°C. The rationale for the

test temperature being higher than the service temperature is two-fold.

First, the actual loading rate in a bridge is substantially slower than the

loading rate in a CVN test. The test temperature is increased by an amount

equal to the shift in transition temperature as the loading rate is decreased

from the impact rate used in CVN tests to the maximum loading rate expected in

bridges. Second, other factors such as service experience, design and fabrication controls, and the results of fatigue tests on similar welded beams also

contribute to the expectation of safe performance by the bridge.

3.4 Summary Comments

Consideration of existing fracture criteria for steel weldments used at

low temperatures does not provide clear guidance for the establishment of

fracture criteria applicable to fixed platforms in the Arctic. Consider, for

example, the range of test temperatures required. The CVN test temperatures

range from 30°C below the minimum service temperature, T , for OCS platforms

to 39°C above T for arctic bridges, with intermediate requirements for ships

and pressure vessels. Clearly, objective criteria are missing.

4. EXPERIMENTAL PROCEDURE

4. 1 Test Material

The test material is a 25.4 mm thick plate of ABS grade EH36 steel, a 350 MPa (51 ksi) yield strength C-Mn steel. The chemical composition is given

in Table 4. The steel is in the normalized condition and has particularly

uniform properties due to sulfide shape control. Thus, material variability

should not have an undue influence on the test results.

4.2 Tensile Tests

Round tensile specimens (6.35 mm diameter and 31.75 mm gage length) with

the tensile axis parallel to the rolling direction were tested at various

temperatures ranging from -196°C to 25°C. All tensile tests were performed in

a screw-driven tensile test machine at a cross-head speed of 0.2 cm/min. The

upper yield point, lower yield stress, ultimate tensile strength, percent

reduction in area, percent elongation and true stress at fracture were measured for each test. All testing and measurement procedures conformed to

the guidelines of ASTM E-8, the standard for tension testing of metals.

- 6 -
The testing machine is equipped with a fixture designed for cryogenic

testing. Various temperatures are attained by an elaborate temperature

control system. The specimen is cooled by reservoirs of liquid nitrogen above

and below the specimen. The specimen grips are equipped with electric heaters

which are connected to controllers. The desired temperature is simply dialed

into the controllers. The combination of heating and cooling the specimen

produces a steady-state temperature.

4.3 Charpy V-Notch Impact Tests

The Charpy- impact transition curve has been established for the normalized steel in the temperature range -196 to 25°C. The procedures of ASTM

E-23, the standard for impact testing, were followed in most cases. Deviations from the standard occurred in some tests. These deviations are noted in

the results and discussion section. For each test the absorbed energy,

percent shear and lateral expansion were measured. The latter two quantities

were measured according to the guidelines which appear in ASTM A-370. The

maximum lateral expansion was measured with a digital micrometer.

The various test temperatures were attained by placing the specimens in a constant temperature bath (liquid nitrogen or a liquid nitrogen - methanol

mixture) for a minimum of ten minutes. Each specimen was then removed from

the bath and tested within five seconds.

4.4 Fracture Toughness Tests

The ductile- to-brittle transition curves have been established as a function of constraint in single-edge notched bend (SENB) specimens (see

Fig. 3). Notch constraint is varied by varying the crack- 1 ength/width ratio

(a/W) and the specimen thickness (B). The specimens are machined with the

notch orientation shown in Fig. 4.

All specimens are fatigue-precracked at room temperature according to the

specifications in the British standard for COD testing [13]. The maximum

allowable stress intensity for fatigue loading, K^-(max), is given by:

K

f (max)= 0.63(jy B

1//2

(1)

- 7 -
where a is the yield strength (in MPa) and B is the specimen thickness

(in mm). The British standard states that the minimum stress intensity,

Kjr(min), be at least 10% of Kp(max), so that:

aK <0.9 K^(max) . (2)

The test matrix is shown in Table 5. For each notch configuration test

temperatures cover the complete fracture mode transition, i.e., cleavage to

ductile tearing.

All tests are performed in displacement control on a 100 kN servohydraul ic test

machine. The displacement rate in all tests is 0.80 mm/min. This corresponds to a loading rate (in the elastic range) for the square section (W = B) specimens of

approximately 9.3 kN/s for shallow notched (a/W = 0.25) specimens and 375 N/s for th

deep notched (a/W = 0.75) specimens. The experimental setup is shown in Fig. 5. Th

mouth-opening displacement and the load-line displacement are measured simultaneous!

during the test by two clip gages. The load-line displacement is measured by the

comparison bar technique developed by Dawes [14]. The load and the two displacement

are recorded on a two-pen X-Y plotter. The test instrumentation is wired to a minicomputer through an anal og-to-digi tal converter. The load, crosshead displacement,

both clip-gage displacements are recorded by the computer at approximately

0.3 s intervals and stored on a magnetic disk. The computer typically

collects and stores around 500 sets of data in a 3 to 5 minute bend test.

Low temperatures are attained by attaching a box to the lower crosshead and

filling it with either liquid nitrogen or an alcohol and dry ice mixture. Intermedi

temperatures (between -196 and -70°C) can be attained by pouring liquid nitrogen int

the box to a level below the specimen. The specimen is cooled by heat transfer thrc

the test fixture and by the vapor coming off of the boiling nitrogen below the

specimen. The actual specimen temperature is measured by a thermocouple implanted i the specimen.

In each test, fracture toughness is measured by the two EPFM parameters, J and

CT0D. Figure 6 shows the notation used for critical values of CT0D. The notation

depends on the nature of the fracture event, i.e. whether the crack extension is

brittle or ductile, and whether or not unstable cleavage is preceded by stable crad

growth. A similar notation is used for reported critical values of the J-integral.
The crack-tip opening displacement (CTOD) is computed from the following

relationship:

K

2 r

n (W" a) Vn

6 =6 +6 = —p + / e p 2a^E r (W-a) , T + a W+ Z (3)

where V = the plastic component of the mouth-opening displacement

Z = knife edge thickness

r = rotational factor.

P

This equation separates CTOD into elastic and plastic components (6 and 6 , respecti vely ) . The British standard [13] suggests that a value of 0.4 be

assumed for r^ in equation (3). However, a more precise value for the rotational factor can be calculated if the plastic components of load-line

displacement and mouth-opening displacement (q and V , respecti vely ) are

P P known [15]:

r

p

“ wri q

2- {(1 - I6W } - (a + z)} (4)

P

In most cases the J-integral for an SENB specimen is estimated from the

following equation

J

= B(W-a)

^

where U is the area under the load/load-line displacement curve. In some

cases other J-estimation procedures are used and the results from the various

procedures are compared. The details of those procedures are given in the

results and discussion.

The method that is utilized to detect the onset of stable crack growth is

the double displacement method. Both the mouth-opening displacements, V, and

the load-line displacements, q, from each test are stored on a magnetic disk.

2 2 A program has been written which computes dq/d V and d q/dV , and plots them as

a function of V. The critical displacement is inferred from these plots.

- 9 -
5. RESULTS AND DISCUSSION

5. 1 Tensile Tests

The tensile properties have been obtained for 10 temperatures ranging

from -196°C to 25°C. These data are shown in Fig. 7 and in Table 6. The data

follow the expected trends with a minimal amount of scatter. Both the yield

strength and the ultimate tensile strength increase markedly with decreasing

temperature. All flow curves exhibited an upper and lower yield point which

is characteristic of low-carbon steels. The yield strength (as plotted in

Fig. 7) was defined as the upper yield point.

5.2 Charpy- Impact Data

The Charpy V-notch impact energy transition curve has been established

and is shown in Fig. 8. The percent shear and lateral expansion were also

measured for each specimen. The Charpy data are summarized in Table 7.

Figure 8 shows that the ductile- to-brittle transition occurs at about -70°C.

Note that the curve is nearly vertical at this temperature. Most specimens

tested at or near this temperature exhibited either fully brittle (low energy)

or fully ductile (high energy) behavior. Only a few specimens exhibited

intermediate behavior. Table 7 shows that the transition from low energy to

high absorbed energy corresponds well with the fracture surface transition (0%

shear to 100% shear) and the transition from small to large lateral expansion.

Because of the high toughness of this material, most of the upper shelf

energy values are invalid according to the ASTM E- 23 standard. The full-scale

energy of the charpy test machine used in this experiment is 358 J. According

to the ASTM standard the energy absorbed by an impact specimen should not

exceed 80% of the full-scale energy of the test machine. Therefore, all

energy values above 268 J do not meet the requirements of ASTM E-23.

Figure 8 illustrates an inherent weakness of material toughness requirements based on Charpy data. For example, if the service temperature is -40°C

the required toughness based on present 0CS platform requirements would be

34 J at -70°C. At this temperature there is a drastic drop in toughness to

well below 34 J while at the service temperature (-40°C) the material is fully

ductile and has a toughness of greater than 250 J.

- 10 -
5.3 Fracture Toughness

The fracture toughness data which have been collected to date are summarized in Figs. 9-14 and Table 8. The critical values of J and CTOD along with

the critical fracture event for each specimen is shown on these figures.

Three critical fracture events have been observed. J and are defined at

c c the occurrence of unstable cleavage without prior stable crack growth. J. and

6. are defined at the onset of stable crack growth. When cleavage is preceded

by stable crack growth J

u and <$ u are measured at the point of instability.

Figures 9 and 10 show that the ductile- to-brittle transition curves shift

approximately 30°C as the specimen thickness increases from 13 to 38 mm. The

thickest specimen produces the highest amount of crack-tip- region constraint;

constraint is the restriction of plastic flow by the surrounding elastic

material. Constraint causes a triaxial stress state and raises the flow

stress near the crack tip. This increase in flow stress tends to promote

cleavage fracture (i.e., shift the transition curve to the right) since it is

easier to reach the fracture stress.

It should be noted that the thickest specimens (B = 38.1 mm) had slightly

deeper cracks than the other two specimen geometries shown in Figs. 9 and 10.

The deeper cracks were necessary in order to prevent the load on the thick

specimens from exceeding the capacity of the test machine. These deeper

cracks, however, probably contributed in part to the shift in the transition

curve to higher temperatures

The influence of crack length on the ductile- to-brittle transition curve

is shown in Figs. 11 and 12. In SENB specimens constraint increases with

crack length because of the plasticity reversal (tension to compression) that

must occur within the remaining ligament. Therefore the transition temperature increases with crack length. Since the crack tip in a shallow-notched

specimen is near a free surface, constraint can be further relaxed by plastic

flow to the free surface. The data for the deep-notched geometry (a/W M3. 75)

behave in a somewhat unusual manner. The curve fitted to these data (Figs. 11

and 12) crosses the curve representing the a/W ^0.5 data. At low values of 6 c and J , where the plastic zone is relatively small, the deep-notched geometry

has more constraint than the a/W ^0.5 geometry. At higher temperatures (and

higher 6 c and J

c values) the plastic zone at fracture is on the order of the

size of the ligament in the deep-notched geometry. The constraint is relaxed

as the plastic zone approaches a free surface.

- 11 -
Figures 13 and 14 show values of 6^. and J., respectively, for all five

specimen geometries. Both lL and 6. are apparently independent of geometry

and 6. is approximately independent of temperature. J. increases slightly

with decreasing temperature because the area under a load-displacement curve

at a constant displacement increases with increasing flow stress. Since the

material used for this study has a high tearing modulus (dJ/da) the point of

incipient crack growth is not well defined, contributing to the scatter seen

in Figs. 13 and 14.

Since increasing strain rate elevates the flow stess, strain rate should

have an effect on the ductile- to-brittle transition. This effect is seen when

the Charpy- impact transition curve in Fig. 8 is compared to the curves in

Figs. 9-12. The ductile- to-brittle transition in the Charpy test (high strain

rate) occurs at about -70°C while all curves in Figs. 9-12 for fracture

toughness tests (low strain rate) lie to the left of this temperature. Also,

note that the transition in the Charpy- impact data is much steeper than the

transition in the fracture toughness data.

Figures 9-12 illustrate a problem facing the field of fracture mechanics.

While the fracture behavior of a steel can be adequately described in terms of

Kj

c on the lower shelf and Jjc and R-curve analysis on the upper shelf, there

is at present no fracture toughness parameter applicable to the ductile- tobrittle transition region. The fracture toughness of a material at a given

temperature in the transition region is dependent on the geometry of the test

specimen

.

5.4 Predicting the Effect of Constraint on the Ductile- to-Brittle Transition

The limit load, P^, or load at net- section yield is related to the yield

stress by the following expression:

L o

y (W-a)

2 B

where L is a dimensionless parameter which is a measure of the constraint due

to the notch. For the material used in this investigation: L = 1.38 for B

=

12.7 mm, a/W = 0.2 and L = 1.52 for B = 38.1 mm, a/W = 0.29.

- 12 -
The quantity Lay is a measure of the average stress in the ligament at

net-section yield. Thus, the ligament stress at fracture, La

c

, can be

computed as follows:

4P W

c B(W-a)

2

( 7 )

where P c is the load at fracture. The quantity Lo

c was found to be relatively

insensitive to temperature and specimen geometry. The average value of Lo

c is

750 MPa for the material used in this study.

A simple model which assumes that Lc

c is independent of temperature and

geometry can be used to predict the effect of constraint on ductile- to-brittle

transition curves. If L is known for a given geometry, the corresponding

value of o (the ligament stress in the absence of triaxiality) can be

computed. The strain required to reach a stress of a

c in a uniaxial tension

test at a given temperature can be infered from the flow curve at that

temperature. CTOD can be infered from ligament strain. It is then possible

to plot 6 c as a function of temperature and L. The above approach was applied to the three shallow-notched geometries.

The resulting curves are plotted in Fig. 15 along with the experimental data.

Although the shape and position of the predicted curves do not agree with

experiment, the relative temperature shifts between geometries was accurately

modeled by this approach.

5.5 Detecting the Onset of Tearing

During a fracture test the load-line displacement (q) is plotted autographical ly against the crack mouth-opening displacement (V). After netsection yield the q versus V plot yields a nearly straight line. However, as the crack grows this line gradually decreases in slope because the center of

rotation moves during crack growth. Figure 16 shows a computer plot of the

first and second derivatives of q with respect to V. The first derivative

plot has proved to be the more useful in detecting crack growth. The behavior

at the far left of this plot represents the transition from elastic to plastic

deformation. During plastic deformation the first derivative remains

relatively constant and then increases sharply. This increase is possibly

13 -
associated with crack-tip blunting. At V = 1.48 mm there is a sharp drop in

the first derivative. This is taken as the point of incipient tearing. At

higher displacements the curve has a zig-zag shape. This is believed to be

associated with successive increments of tearing followed by blunting.

5.6 Estimation of J from CMOD

For deep-notched SENB specimens the J-integral is usually estimated from

the following equation:

2U

T

J " B(W-a)

^

where is the total area under the load versus load-line displacement curve.

The value of J can also be estimated from the load versus mouth-opening

displacement curve using equation

J k

2

, K E

1 B(W-a) (W-a)+a+z' (9)

where is the plastic component of the area under the P/V curve. Equation 9

r is based on the Sumpter and Turner [16] equation, although they approximated

the area by

U

v

P

( 10 )

where is the limit load.

The J-integral has been calculated for a number of tests using

equations 8 and 9 and the agreement is very good. Figure 17 shows a comparison of estimates from the two equations. The data represent a single

specimen at various displacements. The difference between the two J estimates

for this particular specimen range from 0.5% to 1.5%.

The value of r^ used in equation 9 to generate Fig. 17 was measured from

the same specimen using equation 4. The excellent agreement between

equations 8 and 9 as shown in Fig. 17 is evidence that the simple hinge

mechanism is an adequate model to describe the plastic deformation of an SENB

specimen. Figure 17 also shows that the J-integral can be accurately measured

- 14 -
from the mouth-opening displacement if r^ is known. Thus, it is possible to

measure both J and CTOD from an SENB specimen with a single clip gage.

5.7 Relationships Between J and CTOD

Under conditions of small-scale yielding the relationship between J and

CTOD can be estimated by:

J = m ay 6 (11)

where m = 1 for plane stress and m = 2 for plane strain [16]. The value of m for large-scale yielding should lie between 1.0 and 2.0.

The plastic term in equation 9 is similar to the equation for the plastic

CTOD:

6 P

r (W-a)V

-DJ 1_L_

r (W-a)+a+z

P

( 12 )

From equations 9 and 12 one can obtain a simple equation for the ratio of the

plastic J to the plastic CTOD

v 2U W

P,

VyW-a)*

The flow stress can be estimated from the limit-load expression

K

P a flow

(W-a) 2

4W

1.5B

(P>P

L

)

(13)

(14)

Hence m can be calculated:

P

m

P (15)

Equations 13 and 15 have been incorporated into a computer program which

plots J / 6 and m as a function of displacement. A typical set of plots is

r r r

- 15 -
shown in Fig. 18. The ratio J /6 increases with displacement due to work

i r hardening. Also m

p is not constant; it decreases with displacement until it

levels off at a value of 1.27. Apparently the relationship between J and CTOD

is more complicated than has been traditionally assumed. This may have

ramifications with respect to the merits of each as a fracture characterizing

parameter.

5.8 The Eta Factor

A more general form of equation 8 is:

J

= nU

B(W-a) (16)

where n = 2 for deep notched SENB specimens. The above equation can be

divided into elastic and plastic components [16]:

J "e

U e B(W-a)n U + —/ P P . B(W-a)

n U PP

B(W-a) (17)

The elastic eta factor, n , can be derived from the elastic compliance

and stress intensity coefficient. The solid line in Fig. 19 is a plot of n

e versus notch depth [17].

Since J is defined as the negative of the spatial derivative of work, the

plastic eta factor, rip, and the over-all eta factor, n Q

, can be computed as

fol 1 ows

J "o

0

B(W-a)

"I dU

B da

hence

(W-a) dU

% " U da

(18)

(19)

A series of SENB specimens with notch depths ranging from a/W = 0.19 to

0.75 were tested at room temperature. The area under the load versus loadline displacement curve, U, was measured for each specimen at displacements of

1.0, 1.5, and 2,0 mm. These values of U were then plotted versus crack length

- 16 -
and dll/da was measured for various crack lengths by drawing tangents to the U versus a curves and computing the slopes. The over-all eta factor, n > was

then computed as a function of crack length. The results are plotted in

Fig. 19. According to Fig. 19, n Q

is relatively independent of crack length

down to a/W = 0.19. It therefore seems reasonable to use a value of n

= 1.8-

2.0 for all J measurements from SENB specimens with a/W > 0.2.

6. SUMMARY AND CONCLUSIONS

The existing material toughness requirements based on Charpy- impact data

are somewhat inadequate. The goal of this program is to develop a fracture

mechanics rationale for specifying material toughness requirements. This

study, to date, has been devoted to an assessment of the base-metal fracture

properties. Fracture toughness tests on weldments are planned for the

immediate future.

Tensile, Charpy- impact and fracture toughness tests have been performed

on ABS Grade EH36 steel over a wide range of temperatures (-196 to 25°C). The

Charpy-impact transition curve is nearly vertical at approximately -70°C. The

upper-shelf energy of this material is very high (^330 J). Critical values of

J and CTOD were measured for five geometries of SENB specimens. The

ductile- to-brittle transition curves for the SENB tests are much less steep

than the Charpy data; the transition temperatures for the fracture toughness

tests are all lower than that for the Charpy data. Specimen geometry has a drastic effect on fracture toughness in the transition region. Increasing

specimen thickness and/or crack length markedly increases the transition

temperature. Larger specimens produce a greater amount of stress triaxiality

near the crack tip, thus elevating the flow stress locally. A higher flow

stress tends to promote cleavage fracture since it is easier to reach the

fracture stress. Since fracture toughness in the transition region is

geometry dependent, one cannot safely predict the fracture behavior of a large

structure from the toughness of a small, laboratory specimen.

Initial attempts to model the effect of constraint on fracture toughness

have been moderately successful. It is possible to predict the temperature

shifts (due to constraint) in ductile- to-brittle transition curves with the

existing model. Further refinements of the model are needed to accurately

predict the position and shape of transition curves.

17 -
Preliminary data suggest that it is possible to measure both J and CTOD

with a single clip gage which measures the mouth-opening displacement. It was also found that the ratio between J and CTOD is a function of displacement,

yield strength, and work-hardening rate. Direct experimental measurement of

the eta factor, a dimensionless constant used in J calculations, show that it

is approximately equal to 1.85 for displacements of up to 2.0 mm and a/W

ranging from 0.19 to 0.75.

REFERENCES

[1] Broek, D., " Elementary Engineering Fracture Mechanics ,

11 Noordhoff

International Publishing, Leyden, The Netherlands (1974).

[2] Rolfe, S. T. and Barsom, J. M. , " Fracture and Fatigue Control in

Structures, Applications of Fracture Mechanics ," Prentice-Hall, Englewood

Cliffs, New Jersey (1977).

[3] Larson, L. H., Editor, " Advances in El asto-Pl astic Fracture Mechanics ,"

Applied Science Publishers, London (1980).

[4] Burdekin, F. M. and Dawes, M. G., "Practical Use of Yielding and Linear

Elastic Fracture Mechanics With Particular Reference to Pressure

Vessels," Institute of Mechanical Engineers Conference on "Practical

Applications of Fracture Mechanics to Pressure Vessel Technology,"

London, (May 1971).

[5] Anon., "Guidance on Some Methods for the Derivation of Acceptance Levels

for Defects in Fusion Welded Joints," PD6493 : 1980 , British Standards

Institution, London (1980).

[6] Anon., "Requirements for Verifying the Structural Integrity of 0CS

Platforms," U.S. Geological Survey, Reston, VA (Oct. 1979).

[7] Anon., "ASME Boiler and Pressure Vessel Code, Section VIII, Pressure

Vessels," American Society for Mechanical Engineers, New York, NY (1977).

[8] Anon., "Chemical Plant and Petroleum Refining Piping," ANSI B31.3,

American Society for Mechanical Engineers, New York, NY (1977).

[9] McHenry, H. I., "Ship Steel Weldments for Low Temperature Service,"

Welding Journal, 55(5), 1975, pp. 387-393.

[10] Anon., U.S. Coast Guard Marine Engineering Regulations, Subchapter F,

CG-115, U.S. Coast Guard, Washington, DC (July 1970).

[11] Anon., Rules for Building and Classing Steel Vessels, American Bureau of

Shipping, New York, NY (1974).

- 18-
[12] Anon., Standard Specifications for Highway Bridges, American

Association of State Highway and Transportation Officials, Washington,

DC, 12th Edition, (1977).

[13] British Standard 5762: Methods for Crack Opening Displacement (COD)

Testing, 1979, The British Standards Institution, London.

[14] Dawes, M. G., "Displacements and Rotational Factors in SENB

specimens." The Welding Institute Research Bulletin, 7, 1976, p. 185.

[15] Lin, I. H., Dawes, M. G., deWit, R., and Anderson, T. L.,

"Displacements and Rotational Factors in Single Edge Notched Bend

Specimens". Submitted for publication to the International Journal for

Fracture.

[16] Dawes, M. G., "Elastic-Plastic Fracture Toughness Based on the COD and

J-Contour Integral Concepts". In ASTM STP 668, 1979, p. 307.

[17] Sumpter, J. D. G. and Turner, C. E., "A Method for Laboratory

Determination of J ". In ASTM STP 601, 1976, p. 3.

[18] Towers, 0. L., "Stress Intensity Factors, Compliances and Elastic n Factors for Six Test Geometries." Welding Institute Report

7709.01.81/249.1, 1981.

- 19 -
Table

1.

Charpy

V-Notch

Requirements

for

Steel

Plates

in

OCS

Platforms

cj 4-J >1 o3 QJ E E > o E E E O OO O O

O0 1 1 1— >> Ll -(-> 1

-Q Q) QJ QJ 3 3 rC J3

4-> 1

00 rc o 4-> rO CO o o E 2 ^ V

<_> 3 CJ 31 E CJ E

lo

x

O 2 rc QJ - O * E E o 2 E 4-

QJ 3 CO CO XJ 3 E E 3 5 E aj 4-J O CD 3 -»-> rc cn U 4- 4-J

4-> JZ1 E E E E E o rc rc rc

rO 1 f— 1— CO *r- 4-J 00 00 E XJ -1- 5- >> CO OO E 4- aj e

QJ E V

V

V QJ E 00 • QJ •r- CL

CL

O

O 4-J ajx E OJ E 4-

O E cn o C_J CJ 3 to E aj oo •r— •r- E

CD QJ o o o qj aj 3 E 00 o QJ CJ CL

l— +J o o o •r- E +-> QJ -E -3 QJ CL

<E t—1

r—

H

r—

H E rc

x CD 2 4-J CL ra

+-> o O) * E •r— to

00 cl cn aj 00 •I "O 4-J 0J

QJX E Cl E 00 E fO rO E !— CD •»- E O E (C -3 •r- 03 QJ aj CJ CD 4-J E CJ _3 3 -F> 4-> to CD cn •>— rO E E E CJ o >3 aj O 4-J OJ >

1— 1— 1— •r- S +-> +-> 00 to •1— 00 3

>> -E OO 1— E 00 E rO CD E V V V 2 -a CD rc rc QJ 4-J DOW o i +-> •r— a E E

cn CJ CJ CJ •> o CJ •i

+-> rc to -3 3

CD o o o 1— o 4-J QJ 4- oo TD tO

O t4-> o o o CO CO Ol •r— E aj e

rO co CO CO CJ " m E QJ 3 e cn xj

c_> •r- cn +J cn • j— -a -E QJ

+-> E QJ QJ •1

CD •1

— QJ cn E E

•r— *i

E -C OO E to E E -r- CD E -O E 3 O 4-J CJ iE CD TZ 3 4-> QJ >>4J -a e O -Q 1 •i

CJ 4-J i CD E

aj r— 3 rc 3 i QJ cnO E X o +-

E 4—J

rc -EE

U

OJ E -O 1 JE 00 4-> 3 E 4-J E

cn

x s --

—>. 4-J E 00 3 oo rC 3 00 QJ E UD LO <J CL O 00 a 4-J cv j CD 1 1 c—H i—i E 03 -3 00 aj •t— CJ QJ ra 00 > X. ' >— *— E E OO CD o -3 4- E E -r- QJ -—

CO 4- E +-j 4-J •1— E E E E

-Q e O o CO r» \—1 E 3 4-J 4-J OJ E * i— rO C\J oo oo aj e +-> QJ 4- cn OO CJ 4-> 4-J >> E E o QJ CO E O •r— E tV rv cn +-> h- co •!— | E rc OO s- E E e 4- +-> -Q QJ | 1 4-J QJ

CD -

-E rO rc CL El rc o 4-J C/1 CD CL E CJ E (— E CJ > 4-J 3 -E E

Ll! on

— •i- 4-> QJ •l

•i

— QJ CD 4-J QJ CD rc -E E 4- -M -3 > a ’ 1 -E 4-J 2l' r- e ^—. X—v 2 CD O 2 E CD CJ +J 4-J > 3 •P" LO o LO u +-> 3 4-> •1— X> o o X «

1 CO oo cn E 00 OO 00 00 CJ 4- 4- -E E X 3 V V

1—-• E O QJ QJ E CD 4- O 4-J tV

4-J ra CJ 3 +-J rc CD #r

C 3

•1

o LD CJ •r— o -3 -O OO >> E »> cn oo OO CO 00 i

•N +-> 3 4-J o cn E +-> cn 0J CJ +-> oo QJ •r- - 1- CD O E CD X O E o E .— 4-J 4-J _J E •r- E •1 o •1— +-> OO rc •r- (D (O o O 4-J 3 oo O E -Q E E

•1

•f~) oo cn i *r—) -+-> QJ J3 tV

4-J

4-J rc Ol CJ > oo cn

CO

X 4- -LJ s T3 CD E •r— •r- E E E E O 00 o E E QJ • -3 >0-1-

1— CD CO QJ 1— rc 31 E cn 2 E E XJ X 3 4-> CD 4-J E E CD IE

•1 cn

O -Q 00 1 •r~- OO on XJ

O

CD co 00 CO E -r- 4-> E rc rc 4-J E | N CO CO CO E QJ 4-J O OO CD E 00 QJ CD fO

•1 o -Q IE rc CD -O QJ 3 QJ -O _3 DO V 1 A A o QJ E E CL E E oo -M 4-J E +J >, on E E CD -r- E 3 CD a CJ CD -Q E E E f— 1— =3 E -O •1— 4-J E CL 3 OO E O O O CO +-> E rc E to 4-J XS -r-

•1

V •r- rc r- o QJ E 1 XJ 4-J QJ 1 QJ 4-J

4-J CJ E rO CJ -3 CD rc i to 3 rc 4-J CJ CJ CO CD QJ E +j CL E 3 .3 E (O ICE

CD CL CL E i

F 3 o >> cn 3 CD r— E

go oo E +-> IE <n QJ 4-> 2 E 3 4-J E CL

O

QJ CJ 4-J 00 4-> CJ rc o U -E CD CJ ^ 

4-J 3 3 00 3 QJ E 4-J 3 4-J E ^—-N CO E QJ CD QJ E E •1

E cn

CO Cl Q) 4-J E 3 CJ +-J 3 E 4-J 4-J Id Q) E

Q. 5: CJ 00 -I— *f OO CL CJ oo _Q *i

—

•1 E CJ > CJ rC QJ

X o > E +-> •1— E E rc E CL E 00 >, (E LO cn E O QJ -3 QJ o E O E O O CE O 1—1

CD CD Ll XJ 4-> 00 Ll 4- Ll •r~ Ll CL E -—.

co

CO

V

Cl

V E • • 21 >- 3 rc JZi CJ XJ

>- t> E O o •1

>> >> >4 ><

r~H

V E E E E

CO

V •i

—

O

OOO

LO E cn cn cn cn

V o 1—1 QJ aj CD QJ

LlJ 1—1 ll +-> -t-j 4-J 4-J Q >- CO V—

rc rc rc rc

C

D

V—

1—1

SO CJ CJ CJ

C£ ^ 

1—1 1—1 1—

CD i—i i—i 1—1 t—1 *

20
Table

2.

Charpy

V-notch

Requirements

for

Fracture-Critical

Members

in

Arctic

Bridges

X -s .

. X- N. .. CD CD

C-> c_> <—

CJ> CJ o o o o 0 o o o «D- <53- t/5 _Q OJ CM C\J OJ OJ OJ CO CO

to 1— rH I 1 T—

r—

rH r-H 1 1 CD

a> i I 1 1 1 1 8 —

LU

£Z 4-> >. -r

V ^ •—

VD

v ^ • (—

-c q- Li- Ll_ 1—

cn u_ Lu Li— Li_ Ll. Li_ 0 O i—

zs o O o o o O

OO s: o o O O o O O CO CO Od

I— r-H

r—

r-Hr—r-H rH 1 1 LU

o_ T3 -M 4-> -M 4-> 4-> 4-> 4-> +-> a) fX3 03 CO CO CO CO CD CD 1—

-1c. o 

LO LO LO LO LO o LO LO 2: Z5 OJ C\J OJ OJ OJ oo CO CO

cu ^—

*—

x—^ ^ —

s S -V cu —

<3- rH CO 00

oo OO oo OO oo *3" "D- "D"

•a •O TD

CD CL) CU -O "O

c: c E C CU E CU "

-—

<u ^—

CL) E -—CUO E "D

d c +->

d -t—

^—*»

C 4-> 1 i

•r- •r— to

d •r— tod OJ •r~ Ul CU CM QJ CD “1 CD •r— o -O CD 2 o 2

4- =3" 4— r-H CU «D- 4- r-H to -—

OJ OJ TD

— ^ 

-

to to >> 'w- >»

—

o i— cz o c

CL) CU E £ E 4-> CU E •r— 4—

•i

C -£Z E E r— E E i— E 2 E i

E

ZZ

u CD E CD E E CD E E <=3- o cz oo CM U OJ O E ^— C\J O E •1— • 1— o o i

tH o •i— r-H

O •i— =3-

OO

-C r-H 1 i SZ LT) TD r-H e LO "D i—

•1

— \ 1 C to +-> =3- +-> 1— CO aj CD CU LO CD to o o JZ O "TD o JZ o "D t=d- o O LO LO

+->

a 4-> 1

- +->

a +-> i— s- i +->O +-> • s- • cu O) O) CD CU OJ CU OJ CU CM

CL Q. £ CL 2 CL E CL 2 > -CL E CL

V ^ > ^ 

=> ZD ID ID ZD o ZD ZD o

<zz o

•1

—

+J ZI res OJ OJ CO CO CO '=3- =3- 1— cz to o- r-^ CO CO CO i—

r-H i—

oo cn CO tn lo LO LO LO LO LO LO c •<- <c =£ =c <c =£ <c «=c c <c

to

a> a

21

Non

-SI

units

are

listed

first

in

Tables 2&3

in

order

to

be

consistent

with

the

original

regulatory

documents

[7,8,12].
CO 1 O -—

CD c i—

«—

'-D CD 1

x

+ x

—

CD • s: C oo o 1 •i

-O t—

§ E

Cl i— e E

•i

CD

x Q_ 4-> -E CO

4- +J CD f—r-H JZ o o > u C •> < c

CO >o o CO 1 •r— DO x—

r-H 1 (/> E 1

CD c

O) SC o

00 LU •1

oo CO

CD +-> •—

X ..—X sz :> CD lq O 03

CO 00 E pH r-H OJ CL

CD Q. 1

*!

x ^ x ^ V *-

X E E cd s; LU D t—

CD o C\J LO

00 +-> r-H

X

—

r-H i

00 >> oo 03

CD CL

x ^ X >s LS- e D 00 o CD O- CO CD CD r-H C\J C\J -M

-SC i

> «—03 E t_> i

<C 1 o i

CO LO o <+- EC «—

X 1 C\J

oo

SZ<D ^ X

—X E CO 03

<D Q_ Q_

s_ Z si

•l—D r-x LO c

x—

LO

CD LO LO

q;

—

—

.c •r— *i

cd CO CO

+-> CO -X o ! .cc CD CD LO LO 1 i— CD —X r-x. cr> « X > •i— +-> CO 05

OO 00 Q_

V

V CL

>> E SI SI

CL CD in CO E 1— O CO 4-> LO

CD 1—O o LO -E E 1

LO

CD Z3 CD

x

VV

X E

•r— 3 •r—

X

X •1

• C o IT 03 cO CO oo •i

OO 1

Q_ Q_

+-> si si

CDD *o LT) LO (

"D e LO 00 <T> -Q CD -E CD *=3" r-H CD •i— -!-> v | ^T LO A

1— 4- D1 E X. ^

•i

— E o in

CD CD -Q 4-> LO LO 4

CD E E O LO o

CL 4-0 CD OO OO C_D

CO

oo

LD

i-H oo

22
Table

4.

Chemical

Composition

of

the

ABS

Grade

EH36

Steel

(Weight

Percent)

CU

c_>

(T3 -Qo

LOO

ooo

(/) c

QJE

cj

cuQ.

CO

CO2: LlJ COco

lO

to

cu

CO

CO

II CO

co

loooooCO

to

to

O)c

.c:CO

=3o

CD

s-o

CD

S_

LO

CM

II CO

oo

COoo

LO1—

o

o’

s-oX

s_

4-> to

-M

to

cu

r-~. CM

II CO o

CDCO

CM1—

o‘

LO

cu jO

ro

CM LO

LO

r^. • • • O o o

II II II 2 3 2

fC <T3 fO EE

•=d- LOCM

23
Table

6.

Tensile

Properties

of

the

ABS

Grade

EH36

Steel

+-> 0 ro

c: •!- CD CM CD LO OO LO O CD 00 co O t 1 CD +-> S_ u u ra co CO CO *=3

- O 1 1 0 »—

H

CD oc

S- 13

CD "O c

CM CO CO CO f\ r\

Q_ Cl) •>— OC

C

-M O

C -r- CD +-> LO r—

H

CO CD O ro • S- CD CD co 00 CM LO

CD £Z Q_ O

OO 00 00

LO i 1 CO CO LO CM r^- CM CO =3

- O 00 •=3

" 00

<T3 CD D_

C- 21 D

-l-> OO CTl co CM O LO \

1 r-H 0^ O

<a LO OO \ i t—

1

CM •=3- s_

Ll_ CD

S4-> OO

r-H r-H l—

1

r~H 1—

1

CO CO 00 CD co CD 0 C\J CD 00

«—

1

CM CM C\j CM CM OO

*—

1

r-H r—

H

1—

1

1 1 r-H »—

i

CD

•1- ru CO Cl.

CD

CD +-> +-> CD

fD c: £ CD

•i— 4-1 +J OO

co co CM CO OO CD CM LO CD CD o^> 00 CO CO CD '=3

' co i 1 co 00

co CO LO LO i"- co co LO LO

"D rO >— D_

CD 21

>-

CO

S- CO

CD CD 2 i- o +-> _J OO

CO LO LO CD CD 00 O 00 00 I''- 00

CM CM co CM CM r-H co co LO co CO 00 00 CO LO 00 CO LO 00 00

"O

1— <T3 CD Cl.

S- +-> CD C

Q_ •*“ CL O

=> cl1 1 0 co co r^ CO co CD CO co CO *3

- CO r—H r^ r- 1 ^3- CD r-H O

CO CO co CO CO CD *3

- 00

cd

o

00

CD • I— QlolocoocmolO‘=D'—<CD r-~r''-

c\j cm cd co 1—i co co r—. CD co lo <—

i

1—

1

1—t 1—I I I

1—I i—I r-H I I III III

CD I—

CDo ; CDQ. OO

CM OO lo co CO CD

24
Table

7.

Charpy

V-Notch

Impact

Data

for

ABS

Grade

EH36

Steel

c s_ CD 03 O CD

S- -C

<D CO

D_

CO

i_CD +J

03

03CLX

O rO

(D —

=C C

oo

Q.

E

CD

CDO

CDCL CO

Ooo oo

CO oo

o o

o

o

o

oo

lo

lo

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Fracture

of

Steels

31

Figure 1.

The

ducti

le-to-brittle

transition

of

steel

showing

brittle

(linear

elastic),

elastic-plastic,

and

ductile

(fully

plastic)

regions.
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03

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@3

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liaise 1-= E e CO E „ i^i

ss

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32

Figure

2.

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of

Charpy

V-notch

specimens

for

weld

qualification

tests
O M

33

Figure

3.

Single-edge

notched

bend

(SENB)

specimen

used

to

obtain

fracture

toughness

values.
34

Figure

4.

Orientation

of

SENB

specimens

with

respect

to

the

rolling

direction.

Charpy

impact

specimens

were

also

prepared

with

this

orientation.
35

Figure

5.

Photograph

of

the

test

machine

and

three-point

bend

fixture.
5 5 * 2 2 »

05 ^ ©© ©

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is

—

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0)w©

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36
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37

TEMPERATURE
400

seinof ‘A0U3N3 ±0VdlAII AdUVHO

38

Figure

8.

Charpy

impact

transition

curve

for

ABS

grade

EH36

steel.
0.80 i Oon

111

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CC

HIaS

LUH

u

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03 o

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L-3

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39
800 <0 <M

o

s

LU

CC

I-

<

CC

LU

0.

2

LU

H

UJUI-N ‘P ivouiao

40

Figure

10.

Critical

J

for

cleavage

as

a

function

of

temperature

and

specimen

thickness

for

ABS

grade

EH36

steel.

W

=

25.4

mm

(1.0

in).
08*0

ujuj ‘aoio lvouiao

41

Figure

li.

Critical

CTOD

for

cleavage

as

a

function

of

temperature

and

crack

length

for

ABS

grade

EH36

steel.

W

=

25.4

mm

(1.0

in).
800

o

9* LU

CO

<

CO

LU

Q.

LU

H

^.uku-n fi P IVOIliaO

42

Figure

12.

Critical

J

for

cleavage

as

a

function

of

temperature

and

crack

length

for

ABS

grade

EH36

steel.

W

=

25.4

mm

(1.0

in).
0.80

E

o

o

0

1

0

00

1

o

CM

V

o

wuj c ao±o ivoiiiao

43

Figure

13.

Critical

CTOD

for

the

onset

of

tearing

asa

function

of

temperature

and

specimen

geometry

for

ABS

grade

EH36

steel.

W

=

25.4

mm

(1.0 i
700

t

_uiiu-N ‘r 1V0I1SH0

o

UJ

cc

<

CC

UJ

a.

S

yj

44

Figure

14

.

Critical

J

for

the

onset

of

tearing

as

a

function

of

temperature

and

specimen

geometry

for

ABS

grade

EH36

steel.

W

=

25.4

mm

(1.0

in).
0.80

uiuj ‘aoio ivoumo

45

Figure

15.

Comparison

of

predicted

and

experimental

CTOD

values

asa

function

of

temperature

and

specimen

thickness.
SPEC.

C-1,

25

C,

A/W

=

0.2925

LD

<.

!

•<

<

I

* -V ..

Q — W ^

I I I 3AXlVAXd3(j QNC

2: o

i

10 ®

— 1

v_/

Q>O

S g — ®

3 AI1V AId3G IS

46

Figure

16

.

Computer

plots

of

the

first

and

second

derivatives

ofq

with

respect

to

V.

A

sharp

drop

in

the

first

derivative

generally

indicates

tearing.
500

o

uiui-N ‘(seeui b ujojj) p

(D

(0

O

E

47

Figure

17.

Comparison

of

J.

estimates

from

equations

8

and 9

.
in

c\j

O)

CM

<S

II X

<

u

LD

OJ

I o

CJ

LU

CL

00

COIlSVlcDW

ansvid)aoi3/r

w

Q

O

43

Figure

ig.

The

ratio

of

the

plastic

components

ofJ

and

CTOD

and

an

estimate

of

the

plastic

m

factor

as

a

function

of

displacement.
49

Figure

19.

The

overall

eta

factor (n)

asa

function

of

a/W

and

displacement.

The

elastic

eta

factor

cuPve

[17]

is

superimposed

for

comparison.
APPENDIX

Literature Review

This

and recent

literature review covers the basic concepts of fracture mechanics

research pertinent to the present study.

50
CONTENTS

1. LINEAR ELASTIC FRACTURE MECHANICS 53

1.1 The Stress Intensity Factor 53

1.2 Fracture Toughness 54

1.3 Crack Tip Plasticity 55

1.4 Summary Comments 55

2. ELASTIC PLASTIC FRACTURE MECHANICS 56

2.1 The Crack Tip Opening Displacement 56

2.2 J-Integral 58

2.3 Summary Comments 60

3. MICROMECHANISMS OF FRACTURE 61

4. EFFECT OF SPECIMEN SIZE AND CONSTRAINT ON FRACTURE BEHAVIOR 62

5. REFERENCES 64

51
List of Figures Page

Fig. Al. ASTM E- 399 compact specimen for fracture toughness testing. 67

Fig. A2. Three-point bend specimen used to obtain CTOD test results. 68

Fig. A3. Notation for critical CTOD values. 69

Fig. A4. Notation for critical J values. 70

Fig. A5. The three basic micromechanisms of fracture. 71

Fig. A6. Schematic representation of possible flow curves for material 12 near the crack tip, illustrating the effect of flow stress on the micromechanism of fracture, (a) Cleavage, (b) cleavage

preceded by ductile tearing, and (c) stable crack growth.

Fig. A7. Effect of specimen geometry and plastic zone size on fracture 73

toughness. Specimens must lie to the right of the vertical

line to meet ASTM E- 399 specifications.

Fig. A8. Effect of constraint on ductile-to-brittle transition behavior. 74

Fig. A9. Effect of constraint on the micromechanism of fracture. At the 75

temperature represented by the vertical line the low constraint

specimen fails by ductile tearing while the high constraint

specimen fails by cleavage.

Fig. A10. Effect of specimen thickness on the ductile-to-brittle 76

transition behavior.

Fig. All. Effect of notch depth on fracture behavior of SENE specimens. 77

Fig. A12. Effect of type of loading on fracture toughness. 78

Fig. A13. Attempt to match the crack-tip region constraint (CTRC) of a 79

structure with an SENB specimen.
1. LINEAR ELASTIC FRACTURE MECHANICS

The physical basis of fracture mechanics stems from the work of Griffith [1

who demonstrated that the strain energy released upon crack extension is the

driving force for fracture. The strain energy, U, is the work done by load,

P, causing a displacement, A:

( 1 )

where C

= -p- = elastic compliance.

The loss of elastic energy upon crack extension of unit area, A, is

defined as the strain energy release rate, G:

dU

| = dC = dU| dA ' A 2 dA dA 1 P

*

Irwin and Kies [2] used this relationship to determine the fracture resistance

of structural materials, G , by measuring the critical load, P

c

, in a specimen

with a known compliance function, dC/dA.

1 . 1 The Stress Intensity Factor

Irwin [3], who determined the stress distribution near the tip of a crack

located in a linear elastic body, developed a stress analysis basis for

fracture mechanics. The magnitude of the crack tip stress field, a — , was found to be proportional to a single parameter, K, the stress intensity

factor:

a- - = —_f, • (e) = K • f( position) (3) ^ / 2itr

where, r and e are cylindrical position coordinates; r = 0 at the crack tip

and e = 0 in the crack plane. K is a function of the applied stress, a, crack

length, a, and a factor dependent on structural geometry, Y(a):

K = Y(a)a/aT (4)
Tada, Paris and Irwin [4] and Sih [5] have published handbooks of K formulae

for various geometries and boundary conditions.

Irwin [3] demonstrated that the strain energy release rate and stress

intensity factor are related:

K

2 = E'G (5)

where, for plane stress, E' = E, the elastic modulus; and for plane strain,

2

E

1 = E/(l-v ), where v is Poisson's ratio. Thus, it is equivalent to attributing the driving force for fracture to the crack tip stress field, which is

proportional to K, or to the elastic strain energy release rate, G. The

stress intensity, K, is used more commonly than G because K can be computed

for different structural geometries using stress analysis techniques.

1 . 2 Fracture Toughness

Fracture occurs when the crack tip stress field reaches a critical

magnitude, i.e., when K reaches K , the fracture toughness of the material.

K

c is a mechanical property that is a function of temperature, loading rate

and microstructure, much the same as yield strength is; however, K

c is also a function of the extent of crack tip plasticity relative to the other specimen

(or structural) dimensions. If the plasticity is small compared to the

specimen dimensions and the crack size, then K

c approaches a constant minimum

value defined as K^, the plane strain fracture toughness.

Measurements of fracture toughness based on linear elastic theory are limited to the case of plane strain testing in accordance with ASTM E- 399 [6].

K 2 B, a, W-a >_ 2.5 (—) (6)

°ys

where B, a, and W are defined in Fig. Al. The specimens are precracked by

fati gue

cycling to an initial crack length of a/W = 0.5 and subsequently loaded to fail

ure. Kj

c is calculated from the critical load, Pg, the measured crack length,

the specimen dimensions and the specimen calibration function, Y(-n-), as follows

k

q

= ~ y <¥> w B/W

( 7 )

54
K t = Kn if all the conditions of ASTM E399 are met, i.e., precracking proceIC Q

dures, load-displacement record, specimen dimensions, etc.

1.3 Crack Tip Plasticity

Applicability of the linear elastic analysis has been extended to conditions approaching net section yielding by correcting for the zone of plasticity that exists at the crack tip [7]. The idea is that the plastic material

at the crack tip strains without carrying the incremental load; therefore, in

the elastic sense, the crack behaves as if it were slightly longer. The

adjustment is made by adding the radius of the plastic zone, r , to the

apparent crack length:

K = Y(a+r ) a/a+r

y y

r

= 1- (JL)‘

y 2tt a v

( 8 )

(9)

where, a

y is the yield strength at the crack tip. The r^ correction modifies

the crack tip stress field to account for the elastic stress redistribution

due to the localized plasticity.

The extent of crack tip plasticity is influenced by specimen dimensions.

This is particularly true for thickness. As specimen thickness increases, o

y increases from a to /3 o due to through- the- thickness elastic constraint [8].

ys ys 3 The maximum value of is reached when the plastic zone size is limited to

about 5% of the thickness. Thus, in a given material, the plastic zone size

as computed by Eq. (9) can vary with thickness by a factor of three -- leading

to a strong dependence of K

c on thickness.

1.4 Summary Comments

The crack tip stress field is the driving force for fracture and the

magnitude of this stress field is proportional to the stress intensity factor,

K. K is a function of crack size, applied stress and structural geometry and

can be computed using stress analysis methods. The resistance to fracture is

a material property defined as the fracture toughness, K

c

; fracture occurs

when K = K . Thus, for a given structural geometry and material, critical

crack sizes can be determined as a function of applied stress and vice-versa:
( 10 )

K 2 a

c = •

The approach is applicable to conditions of localized crack tip plasticity,

i.e., where a £ .8a or r <_ 0.3a, when the r correction is used to account

y ^ y y

for crack tip plasticity [9].

2. ELASTIC-PLASTIC FRACTURE MECHANICS

Many of the high toughness structural materials used for arctic structures

undergo extensive plastic deformation prior to fracture. Thus, the concepts

of linear elastic fracture mechanics must be extended to account for elasticplastic behavior. The two approaches of interest in the present investigation

are the crack tip opening displacement (CTOD) and the J integral.

2.1 The Crack Tip Opening Displacement

The CTOD concept is a crack tip strain criterion for fracture stemming

primarily from the work of Wells [10,11]. Wells' hypothesis was that the

opening of the crack faces was strictly related to the crack extension force,

G, and the stress intensity factor, K, under locally plastic conditions.

Wells used an energy balance argument to derive an expression for CTOD [11]:

<S Eo

y

(ID

where 6 is the CTOD.

In 1965 Burdekin and Stone [12] illustrated how the CTOD concept could be

used to extend the capability of conventional fracture mechanics to the

el astic- pi astic case. A Dugdale strip-yield model [13] was used to develop

the following equation for plane stress Mode I crack-tip displacements under

monotonic loading.

6 = —p 1 n sec (5 —) (12)

" E 2 °y

A fracture toughness test method based on the CTOD concept was developed

by investigations at the Welding Institute [14-16]. The CTOD test specimen

contained a fatigue-precracked notch and was loaded in three-point bending to

56
fracture. The critical CTOD was obtained from the load versus clip-gage

displacement record. The clip-gage displacement, V , was measured across the

notch mouth as shown in Fig. A2. The V was converted to CTOD using the

9

following relationship [14].

6 =

1 + ( a+z

r(W-a

)

(13)

where z = knife edge thickness

a = crack length

W = specimen width

r = rotational factor

This equation was derived assuming that the crack faces opened by a simple

hinge mechanism about a center of rotation. The value of r was assumed to be

0.5 in early work.

The current British standard for CTOD testing [17] specifies that the

CTOD be calculated from the following relationship.

i/2 0.4(W-a)V

&0 = + P (14)

2a E

1 0.4W + 0.6a + z y

where the first term is the elastic component of CTOD, 6 Qi and the second term

is the plastic component, 6 . The plastic component of the clip-gage displaceP ment, V , is estimated by constructing a line, parallel to the elastic-loading

P line, from the critical point on the load-displacement curve to zero load.

The elastic component of the clip gage displacement is then subtracted from

the total displacement. The stress intensity factor for the elastic CTOD

calculation is obtained from the following relationship.

K

= YP

BW

1/2 (15)

where P is the applied load and Y is the stress intensity coefficient. Y is a function of specimen geometry and can be obtained for a given crack length to

width ratio (a/W) from the test standard [17].

57
A more detailed description of the CTOD test method is given in the body

of this report. The type of critical CTOD reported depends on the nature of

the observed fracture event. This is illustrated schematically in Fig. A3 for

a steel which undergoes a ductile-to-brittle transition. At low temperatures

the steel fails by cleavage and 6 is measured experimentally. As the test

temperature increases cleavage becomes less favorable and the fracture tough- ness increases. Eventually the fracture mode changes to microvoid coalescence

and the crack grows in a stable manner. 6. is defined as the value of CTOD at

the onset of tearing. At temperatures slightly above the fracture mode change,

stable tearing can be followed by unstable cleavage. When this occurs, 6 u is

measured at the instability point. On the upper shelf of toughness, the steel

reaches a point of plastic collapse when the work-hardening cannot keep pace

with the decrease in ligament area caused by stable crack growth. 6 m

is then

measured at the point of maximum load in a bend test. 6^ is the total CTOD,

i.e. 6. plus the portion of CTOD associated with tearing. A plot of 6^ versus

Aa, is the resistance (R) curve, , a measure of a material's resistance to

crack growth. Figure A3 also shows that 6. is not a strong function of

temperature, but it is a function of the inclusion content.

2.2 J- Integral

The J-integral is a character!

- zation developed by Rice [18] of the

elastic-plastic field in the vicinity of the crack tip. The use of the

J-integral as a fracture criterion was suggested and experimentally demonstrated by Begley and Landes [19]. J is defined as the line integral:

J = f

T [wdy - f (f^)ds] (16)

where, r is any contour surrounding the crack tip, w is the strain-energy density,

f is the stress vector normal to r,

u is the displacement vector, and

s is the arc length along r. The J-integral is path independent for linear and non-linear elastic materials

[18] and nearly so for most structural materials under monotonic loading

conditions [20]. Thus, J can be computed using numerical methods by analyzing

58
along a contour away from the crack tip, i.e., in a region where the analysis

methods are quite accurate. This relieves the uncertainties of the crack tip

region -- a problem that seriously limits the usefulness of the CTOD method.

An equivalent interpretation is that J is equal to the change of potential

energy U, upon an increment of crack extension of unit area, A:

J --$- (17)

For the linear elastic case the potential energy equals the strain energy and,

therefore, Eq. (17) is the same as Eq. (2) and J = G. Thus, J appears to be a logical extension of linear elastic fracture mechanics into the elastic-plastic

range. Due to the irreversibil ity of plastic deformation, the energy interpretation of the J- integral does not apply to the process of crack extension in

elastic-plastic materials -- as G is for elastic materials. J is simply an analytically convenient, measurable parameter that is a characteristic of the

elastic-plastic field at the crack tip. The contribution of Begley and

Landes [19] was to demonstrate that crack initiation under el astic- pi astic

conditions occurs at a characteristic value of J, called J T , that is related

lc

to Kj

c in the same way G is related to K in Eq. (5). Thus, J integral methods

can be used to determine Kj

c in specimens significantly smaller than the size

requirements of Eq. (6).

Methods of measuring Jj

c are covered by ASTM standard E-813. A deeply

notched specimen of the compact tension or single-edge-notch-bend design is

precracked to a/W = 0.6. The specimen is loaded incremental ly to a series of

J-levels. After each increment of loading, J and the crack extension, Aa, are measured. The load-displacement curve is recorded on an X-Y plotter. Displacement is measured along the load line. J is calculated from the loaddisplacement record and specimen dimensions using:

J

i = iA

bB» (18)

where A is the area under the load displacement record, b is the ligament

(W-a), and n is a dimensionless parameter dependent on specimen geometry. Aa

is measured by either the crack marking technique or the unloading compliance

technique described in ASTM E-813. The results of a test series are plotted

59
as J versus Aa, and the Jj is defined as the extrapolation of the best-fit curve to the blunting line defined by J = 2a v Aa. K T is related to J T as

i ic ic

fol 1 ows

K Ic

(J) = ^

(19)

A more detailed description of the J-integral test method is given in the body

of this report.

The Jj

c test method is limited to measurement of the onset of ductile

tearing. Thus, it is not applicable to the measurement of fracture toughness

when cleavage is the micromode of fracture. Recently Dawes [26] has proposed

the adoption of a critical J notation equivalent to the CTOD notation of

Fig. A3. This notation, shown in Fig. A4, would make the J concept applicable

to all micromechanisms of fracture. This proposed notation is used in the

present study.

2.3 Summary Comments

The CTOD and J-integral concepts are efforts to define single parameter

characterizations of the fracture process that are applicable to linear

elastic and general yielding fracture. In the CTOD concept, attention is

focused at the crack tip; fracture theoretical ly occurs when a critical

displacement develops at the crack tip. The J-integral concept examines the

stress- strain conditions along an arbitrary contour away from the crack tip;

fracture theoretical ly occurs when the potential energy available for crack

extension reaches a critical value defined by J. For the linear elastic case,

6 and J T are consistent with the linear elastic fracture criteria:

c Ic

G

Ic E

1 ( 20 )

There has been a great deal of research in recent years on ductile

fracture. Models have been developed to predict fracture by tearing instability using J-integral R-curves and to predict fracture by plastic collapse

using limit load expressions [21]. This work is not reviewed here because it

is not directly relevant to fracture of structural steels at low temperatures.

60
3. MICROMECHANISMS OF FRACTURE

Fracture usually occurs by one of three mechanisms: ductile tearing,

brittle cleavage, or intergranular cracking. The three mechanisms, shown

schematically in Fig. A5, are competitive. That is, the fracture mode will

depend on which condition is reached first: the critical strain for ductile

tearing, the critical stress for cleavage, or the stress that exceeds the

cohesive strength of the grain boundary. The conditions for ductile tearing

and brittle cleavage are represented in Fig. A6. Note that fracture may

initiate by ductile tearing and, as the stress increases, convert to cleavage.

Fracture toughness is strongly influenced by the micromode of fracture, and

thus it is important in a study of the ductile-to-brittle transition to

identify the governing fracture mode.

To be suitable for OCS platforms, structural alloys should fracture in a ductile manner at the minimum service temperature. Ductile tearing, also

called microvoid coalescence, is caused by the formation and growth of the

voids that eventually comprise the fracture surface. Thus, the toughness of

ductile metals is related to the factors that influence the nucleation and

growth of voids. Voids nucleate most readily at second-phase particles, such

as inclusions and precipitates (1 to 10 ym in size), as a result of interfacial

separation, fracture of the particle, or matrix separation caused by strain

concentration near the particle. Voids grow and coalesce by ductile tearing

of the matrix. Ductile tearing resistance is a function of the strength and

ductility of the matrix. As matrix strength increases, less energy is dissipated by plastic deformation during tearing and toughness is reduced. Increased

matrix strength also tends to activate additional void nucleation sites.

Consequently, there is usually an inverse relationship between fracture

toughness and yield strength when the fracture mode is ductile tearing.

Brittle fracture requires less energy for surface formation than ductile

fracture. Lower energy fracture modes include cleavage and intercrystall ine

(grain boundary) fracture. Cleavage is a fracture mode in which material

separation proceeds along preferred crystal lographic planes without sizable

plastic deformation prior to fracture. Metals subject to cleavage, such as structural steels, usually have a large increase in yield strength as temperature is decreased. Cleavage occurs when the cleavage fracture stress is

reached before the strain required for void formation is exceeded. Intercrystalline fracture occurs when the cohesive strength of the grain boundary is

61
exceeded before cleavage or ductile fracture occurs. As matrix strength

increases with decreasing temperature, intergranul ar failure may occur more readily in a susceptible alloy. However, as a class, the structural steels

that are candidates for arctic construction are not susceptible to intergranular fracture.

4. EFFECT OF SPECIMEN SIZE AND CONSTRAINT ON FRACTURE BEHAVIOR

Specimen geometry can have a substantial effect on fracture toughness as illustrated in Fig. A7. As thickness, crack length, and uncracked ligament

length become large compared to the plastic zone the fracture toughness

reaches a minimum value defined by K^. The ASTM E- 399 standard is designed

to measure this lower bound fracture toughness. The behavior in Fig. A7 is

due to the fact that large specimens promote a maximum degree of triaxiality

at the crack tip, thereby increasing the yield stress. This increase in yield

stress tends to promote fracture since it is easier to reach the fracture

stress. Smaller, less constrained specimens tend to lose triaxiality so that

a greater amount of plastic deformation must occur before the fracture stress

can be reached in the plastic zone. The amount of constraint in the region of the crack tip will depend on the yield strength of the material and the geometry of the fracture specimen.

The specimen geometry consequently has an effect on the ductile-to-brittle

transition behavior of steels. This is illustrated schematically in Fig. A8.

The lower-bound LEFM curve is shown along with four hypothetical specimen

geometries with varying degrees of constraint. The lowest constraint geometry,

represented by the curve farthest to the left, has the lowest transition

temperature and has the highest toughness at a constant temperature. The

vertical line in Fig. A8, drawn at an arbitrary temperature in the transition

region reveals that a wide range of toughness values could be obtained for a material at a given test temperature, depending on the constraint of the

fracture specimen.

Figure A9 shows how varying degrees of constraint can lead to different

micromodes of fracture at a constant temperature. This represents a potentially dangerous situation when using fracture toughness data from small

specimens to design large structures. At temperatures where ductile tearing

is the micromode of fracture, it is permissible to predict the fracture

toughness of large structures from small specimen data because J. and 6^ are

62
not a strong function of geometry [22,23]. However, at temperatures where

cleavage may be the micromode of fracture, a crack in a small fracture specimen

may grow in a stable and ductile manner while a large structure could fail

catastrophical ly by cleavage and have much lower toughness than predicted in

the small scale test.

There is a very limited amount of data in the literature which shows the

effect of specimen geometry on fracture toughness. An added handicap is that

many of the authors do not specify the micromode(s) of fracture which they

observe. Since it is impossible at this time to describe quantitatively the

effect of specimen size on constraint the following section will offer a purely qualitative description based on theory and a limited amount of experimental data.

A number of investigators [23-28] have studied the effect of specimen

thickness on the fracture toughness of various materials. Figure A10 illustrates the effect of thickness on the ductile-to-brittle curves of ferritic

steels. Note that an increase in thickness from B to 2B has a greater effect

than increasing thickness from 2B to 3B. This is because increases in thickness increase constraint at a decreasing rate until a saturation thickness

(plane strain) is reached.

It should be noted that not all investigators attribute the decrease in

toughness with specimen size to constraint. Although the average toughness of

the smaller size specimens is higher than the Kj

c value, the lowest toughness

values for the smaller specimens can be lower than the large specimen Kj

value. This scatter phenomena can be explained by the statistical effect of

inhomogeneity along the crack front [29]. This inhomogeneity will cause greater scatter in small specimen results because the volume of material

sampled is less. Landes and Begley [29] maintain that this statistical effect

alone is sufficient to explain the size effect on fracture toughness. They

applied a Weibull [30] statistical distribution to predict Kj

c from scattered

small specimen results. However, Lin [31] believes that this approach gives a too conservative estimate of Kj

c

. Lin points out that the model of Landes and

Begley based on the Weibull distribution erroneously predicts that

approaches zero as the size of samples approaches infinity. Also, Lin used

63
the Ritchie-Knott-Rice model for cleavage fracture [32] to show that fracture

toughness increases from plane strain to plane stress conditions.

It is likely that the size effect on fracture toughness is caused both by

constraint and statistical effects. It is difficult, however, to separate out

the statistical effects in order to quantitatively study constraint. The

experimental scatter in the ductile-to-brittle transition region is often so great that it masks the constraint effects.

The relative crack length and ligament length can also affect crack tip

constraint. Figures A7 and All show schematically the consensus results of a number of investigations [22-28, 33-35] of the effect of crack length to width

ratio. At a constant temperature the fracture toughness is relatively high

for short cracks but decreases rapidly with crack length until a lower plateau

is reached (Fig. A7). For very deep cracks (a/W > .75) the toughness

increases rapidly with crack length. This phenomena can be explained as follows [36]. When the crack tip is near the center of the specimen the

constraint is maximized because a maximum amount of material surrounds the

crack tip. When the crack tip is near a free surface, as in a very deep or shallow crack, the constraint can be relaxed by plastic deformation at the

free surface.

In addition to specimen geometry, the type of loading can affect cracktip constraint. Bending produces more constraint than tensile loading because

of the compressive stress below the neutral axis of a bend specimen. The

corresponding effect of loading type on the fracture behavior of steels is

illustrated in Fig. A12. The Welding Institute in Great Britan utilizes this

effect when assessing the fitness for service of structures containing cracks

[36]. Figure A13 shows a structure with a hypothetical defect along with the

corresponding SENB specimen. The Welding Institute uses a square-section

specimen with the depth, W, matching the plate thickness and the crack length,

a, matching that of the structure. It is then assumed that the increased

constraint in the SENB specimen due to bending will offset the loss of con- straint due to the lack of plane strain in the test specimen. This assumption

has not been satisfactori ly verified.

5. REFERENCES

[1] A. A. Griffith, Phil. Trans. Roy. Soc. 221A, 1920, p. 163.

[2] G. R. Irwin and J. A. Kies, Welding J., 33, 1954, p. 193s.

64
[3] G. R. Irwin, Trans. ASTM, J. Appl . Mech., 24, 1957, p. 361.

[4] H. Tada, P. C. Paris, and G. R. Irwin, The Stress Analysis of Cracks

Handbook, Del Research Corp., Hellertown, Pennsylvania (1973).

[5] G. C. Sih, Handbook of Stress Intensity Factors for Researchers and

Engineers, Institute of Fracture an Solid Mechanics, Lehigh University,

Bethlehem, Pennsylvania (1973).

[6] 1974 Annual Book of ASTM Standards, Part 10 ASTM, Philadelphia (1974),

p. 432.

[7] G. R. Irwin, in: Proceedings Seventh Sagamore Ordinance Materials

Research Conference, Syracuse University Press, Syracuse, New York

(1960), p. IV-63.

[8] G. R. Irwin, Trans. ASME, J. Basic Eng. 82, 1960, p. 417.

[9] P. C. Paris and G. C. Sih, in: Fracture Toughness Testing and Its

Applications, ASTM STP 381, 1965, p. 30.

[10] A. A. Wells, Brit. Welding J., 12, 1965, p. 2.

[11] A. A. Wells, Brit. Welding J., 10 (11), 1963, p. 563.

[12] F. M. Burdekin and D. W. Stone, J. Strain Anal., 1966.

[13] D. S. Dugdale, J. Mech. Phys. Solids, 8 (2), 1960, p. 100.

[14] A. A. Wells, Congress of Applied Mechanics, Calgary, 1971, p. 57.

[15] F. M. Burdekin and M. G. Dawes, Proc. Mech. Eng. Conf., London,

May 1971.

[16] M. G. Dawes, Advances in Elasto-Plastic Fracture Mechanics, edited by

E. H. Larsson, Applied Science Publishers, London, 1980, p. 279.

[17] British Standard 5762: Methods for Crack Opening Displacement (COD)

Testing, 1979, The British Standards Institution, London.

[18] J. R. Rice, J. Appl. Mech., 35 (2), 1968, p. 379.

[19] J. A. Begley and J. D. Landes, ASTM STP 514, 1972, p. 1. [20] D. J. Hayes, Ph.D. Thesis, University of London, 1970.

L 2 1 ] J. G. Blavel, Advances in Elasto-Plastic Fracture Mechanics, edited by

E. H. Larsson, Applied Science Publishers, London, 1980, p. 65.

[22] M. G. Dawes, ASTM STP 668, 1979, p. 307.

[23] A. B. Stuber, Ph.D. Thesis, University of Kansas, 1980.

[24] R. 0. Ritchie, R. F. Smith, and J. F. Knott, Metal Science, 9, 1975,

p. 485.

[25] A. Penelon, M. N. Bassim, and J. M. Dorlot, ASTM STP 677, 1979, p. 449.

[26] I. Milne, Mat. Sci . Eng., 30, 1977, p. 243.

[27] G. G. Chell and G. M. Spink, Eng. Frac. Mech., 9, 1977, p. 101.

65
[28] I. Milne and P. J. Worthington, Mat. Sci. Eng., 26, 1976, p. 185.

[29] J . D. Landes, and J. A. Begley, ASTM STP 632, 1978, p. 57.

[30] W. Weibull, J. Appl . Mech., 18, 1953.

[31] I. H. Lin, Unpublished work.

[32] R. 0. Ritchie, J. F. Knott, and J. R. Rice, J. Mech. Phys. Solids, 21,

1973, p. 395.

[33] G. G. Chell and A. Davidson, Mat. Sci. Eng., 24, 1976, p. 45.

[34] G. G. Chell and R. S. Gates, Int. J. Fract., 14 (2), 1978, p. 223.

[35] P. M. S. T. deCastro, J. Spurrier, and P. Hancock, ASTM STP 677, 1979,

p. 486.

[36] M. G. Dawes, Private communication.

66
67

Fig.

Al.

ASTM

E-399

compact

specimen

for

fracture

toughness

testing.
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I I I I

Ductile Tearing

t t t

III!

Cleavage

t t t t

Intergranular

III*

Cracking

Fig. A5. The three basic micromechanisms of fracture.

71
Cleavage

72

Schematic

representation

of

possible

flow

curves

for

material

near

the

crack

tip,

illustrating

the

effect

of

flow

stress

on

the

micromechanism

of

fracture,

(a)

Cleavage,

(b)

cleavage

preceded

by

ductile

tearing,

and

(c)

stable

crack

growth.
Small

scale

yielding

CM

r-- <C

CD

73

Effect

of

specimen

geometry

and

plastic

zone

size

on

fracture

toughness.

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the

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A9.

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the

micromechanism

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the

temperature

represented

by

the

vertical

line

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constraint

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fails

by

ductile

tearing

while

the

high

constraint

specimen

fails

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79

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NBS-114A (REV. 2-80

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

1. PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No. 3. Publ i cation D ate

SHEET (See instructions) NBSIR 83-1680 December 1982

4. TITLE AND SUBTITLE

Interim Progress Report:

FRACTURE TOUGHNESS OF STEEL WELDMENTS FOR ARCTIC STRUCTURES

5. AUTHOR(S)

T. L. Anderson and H. I. McHenry

6. PERFORMING ORGANIZATION (If joint or other than NBS, see instructions) 7. Contract/Grant No. NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street, City. State. ZIP)

U.S. Department of Interior

Minerals Management Service

12203 Sunrise Valley Drive

Reston, VA 22091

10.

SUPPLEMENTARY NOTES

Document describes a computer program; SF-185, FlPS Software Summary, is attached.

11.

ABSTRACT (A 200-word or less factual summary of most si gnificant information. If document includes a significant

bi bl iography or literature survey, mention it here) This report summarizes the progress in the development of fracture criteria for steel

weldments in arctic structures. Tensile, Charpy-impact , and fracture toughness

properties have been measured as a function of temperature for a 25.4 mm (1 in) thick

plate of normalized steel. Fracture toughness tests were performed on five geometries of single-edge notched bend (SENB) specimens. Critical values of the J-int^;raL

and the crack-tip opening displacement (CTOD) were computed and plotted versus

temperature. The ductile-to-brittle transition temperature increased with increasing

specimen thickness, and crack length. The effect of specimen geometry on fracture

toughness is attributed to changes in crack-tip region constraint with geometry.

Initial attempts to model this behavior have been moderately succesful.

Various aspects of the SENB fracture toughness test are being examined. Preliminary

results indicate that the J-integral can be accurately measured with either load-line

displacement or mouth-opening displacement measurements. It may therefore be possible to measure both J and CTOD with a single clip gage. The relationship between

J and CTOD has been investigated. The ratio of J to CTOD is a function of yield

strength, displacement and work-hardening rate. The eta factor, which is a dimensionless constant used to relate the J-integral to energy absorbed by the specimen, was

found to be independent of crack length for SENB specimens with a/W ranging from

0.19 to 0.75.

12.

KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semi colon s) arctic structures; crack-tip opening displacement; ductile-to-brittle transition;

elastic-plastic fracture; J-integral; mechanical properties; structural steels;

toughness

13.

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Items 4, 5 — Complete as shown on manuscript. When NBS-II4A is resubmitted along w ith NBS -266

,

following publication of non-NBS media papers, these items must agree with published paper.

Item 6 — If not NBS, blank out and enter G ran tee/ Con tractor name and address, or if performed

jointly, show both.

Item 7 — Complete when applicable.

Item 8 — Enter "Interim,” "Final,” or period covered.

Item 9 — Enter all sponsors’ names and addresses. Include NBS if also a sponsor.

Item 10 — Enter other relevant information, i.e., related or superseded documents. Also used by

Publications Office for Library of Congress catalog number, and entry of non-N BS media

citation upon receipt of Form N BS-266 from author. Check block if appropriate and attach SF 1 85.

Item 11, 12 — Prepare abstract and key words with special care. These are published separately by NBS, NTIS, and other bibliographic services, and are vital elements in guiding readers to your paper. The key words will be used as entries in a subject index. See NBS Communications Manual

for additional guidance.

Item 13 — Indicate "Unlimited” — for open-literature documents cleared under NBS editorial procedures,

or “For official distribution. Do not release to NTIS” — for limited, restricted, or need-to- know material (See Communications Manual). Publications Office will mark appropriate

"order” box and complete Stock Number when known.

Items 14, 15 —Leave blank. To be completed by Publications Office or call Printing and Duplicating for NBSIR’s.

NBS-1 1 4 A (REV. 2-80)