REFERENCE

NBS

PUBLICATIONS

NBSIR 85-3020

NAT L INST. OF STAND & TECH

AlllDb 31A3 c

i DUCTILE-TO-BRITTLE TRANSITION

IN STEEL WELDMENTS FOR ARCTIC

STRUCTURES

F. Zia-Ebrahimi

National Bureau of Standards

U.S. Department of Commerce

Boulder, Colorado 80303

April 1985

• U56

85-31)20

1985

NATIONAL BUREAU Of standards

library

NBSIR 85-3020

DUCTILE-TO-BRITTLE TRANSITION

IN STEEL WELDMENTS FOR ARCTIC

STRUCTURES

F. Zia-Ebrahimi

Fracture and Deformation Division

Center for Materials Science

National Bureau of Standards

U.S. Department of Commerce

Boulder, Colorado 80303

April 1985

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

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

CONTENTS

Page

1 . Introduction 3

2. Material and Experimental

Procedures 7

2.1 Test Material 7 2.2 Tensile Testing 7 2.3 Charpy V-Notch Testing 8

2.4 COD Testing 8

2.5 Microscopy 10

3. Results 10

3.1 Optical Microscopy 10

3.2 Tensile Properties 11

3.3 Impact Toughness 11

3.4 Fracture Toughness 12

3.5 Fractography 12

3.5.1 SENB Specimens 12

3.5.2 CVN Specimens 16

4. Comparison With Base Metal Properties 17

4.1 Tensile Properties 17

4.2 Impact Toughness 17

4.3 Fracture Toughness 18

5. Discussion 19

6. Summary and Conclusions 25

7. References 27

iii

DUCTILE-TO-BRITTLE TRANSITION IN STEEL WELDMENTS FOR ARCTIC STRUCTURES

F. Zia-Ebrahimi

Fracture and Deformation Division

National Bureau of Standards

Boulder, Colorado 80303

This report summarizes the work performed in support of the development

of fracture criteria for steel weldments in arctic structures. The ductileto-brittle transition behavior of a shielded metal-arc weld, typical of steel

weldments in arctic structures, has been studied. Fracture toughness, Charpy

V-notch impact energy, and tensile properties have been measured as a function

of temperature throughout the ductile-to-brittle transition range. The effect

of geometric dimensions on fracture toughness has been studied for three

geometries of single-edge-notch-bend (SENB) specimens. The fracture surfaces

of broken specimens have been characterized by scanning electron microscopy

(SEM) . The microstructure of the multiple-pass weldment has been studied by

optical microscopy. The mechanical properties of the steel weldment have been

compared to the base metal, an ABS grade EH36 steel in normalized condition.

The results indicate that the resistance to brittle fracture is the same

in both the base metal and the weld metal; however, the weld metal has a much

lower resistance to ductile fracture than the base metal does. The high

ductile-to-brittle transition temperature of the weld metal is attributed to

its high tensile strength and low resistance to ductile fracture.

The ductile-to-brittle transition temperature increased by 25°C with

increasing specimen thickness (from 12.7 mm to 25.4 mm), an effect that can be

1
attributed to both the constraint and statistical effects. The slight

increase in ductile-to-brittle transition temperature with increasing crack

length/specimen width ratio is discussed on the basis of constraint effects.

The SEM fractography revealed a shear zone preceding the unstable brittle

fracture in SENB specimens tested in the transition temperature regime. This

finding shows that the fracture initiation process in these specimens is

controlled by ductile fracture rather than brittle fracture, although no

macroscopic slow crack growth was observed.

Key words: crack initiation; fracture toughness; fracture surface; microstructure; 3-point band tests; transition temperature; weldments

2
1. Introduction

The reliability and overall safety of a structure depend on all phases of

design, fabrication, quality control, and inspection, by which a fractureresistant structure is constructed. Fracture mechanics is a tool to translate

the results of fracture toughness testing into engineering design. Although,

many recently developed structural codes and specifications that have material

toughness requirements are based on fracture mechanics principles, the toughness tests specified for material purchase or quality control are in terms of

simple tests, such as the Charpy V-notch (CVN) test. The inadequacy of CVN

testing for evaluation of the fracture toughness of steel weldments in arctic

structures and the need for development of more satisfactory fracture criteria

have been discussed elsewhere [1,2].

Since the results of fracture toughness testing are obtained in the

presence of a notch or some other form of stress raiser, knowledge of the

precise magnitudes and distribution of stresses and strains near the stress

concentrator is essential. To correlate the results of fracture toughness

testing with a parameter that represents a material property, a fracture

criterion should be established. Theoretically, the problem of stresses and

strains ahead of a crack has been solved for small-scale yielding and fully

plastic cases for either plane-strain or plane-stress conditions. No theoretical solution is available for situations where the amount of plastic deformation at the crack tip is between small-scale yielding and general yielding,

and where the stress state is a mixed condition of plane-strain and planestress conditions. As far as the fracture criteria for various modes of

fracture are concerned, it is commonly assumed that brittle fracture occurs

3
when a critical fracture stress is achieved ahead of the crack tip (stresscontrolled fracture) and that ductile fracture occurs when a critical fracture

strain is obtained ahead of the crack tip (strain-controlled fracture) [3],

In general, unstable fracture occurs when the driving force for crack

extension exceeds the material's resistance to fracture. Under small-scaleyielding conditions, the driving force for extension of a stationary crack can

be related to the crack-tip stress intensity factor, K. Unstable fracture

occurs when the stress intensity factor reaches its critical value, K^, which

is a measure of the fracture toughness of the material. Linear-elasticfracture mechanics (LEFM) provides the relationship between K and macroscopically measurable parameters, namely applied stress and crack length, for

plane-strain and plane-stress deformation conditions at the crack tip.

Plane-strain fracture toughness, K^, testing (ASTM E-399) has been designed

to measure the toughness of a material under small-scale-yielding and planestrain conditions where principles of LEFM are valid. These conditions are

achieved through specimen size restrictions in the test method. The

fracture criterion defined by plane-strain fracture toughness testing is

geometry independent and, therefore, the K value is a material property.

However, plane-strain fracture toughness is a conservative measure of fracture

toughness of a material and for laboratory-size specimens the applicability of

the Kt test method is limited to relatively brittle structural materials, to

Ic

low testing temperatures that are below normal service temperature, or to very

high rates of loading.

In the absence of sufficient constraint at the crack tip, stable or slow

crack growth may occur in relatively high-toughness materials. The R-curve

testing method (ASTM E-561) has been designed to characterize the resistance

of cracked thin metal sheets to stable crack growth under small-scale

4
yielding. Under limiting conditions of plane stress or plane strain, the

increase in crack length, Aa, has a unique functional relationship to the

stress intensity factor, K. If fracture, stable or unstable, occurs under a

mixed condition of plane strain and plane stress, the stress state ahead of

the crack is complex and a single parameter, K, cannot characterize the stress

distribution. The complexity increases when crack extension occurs under

large-scale yielding, i.e., the size of the plastic zone at the crack tip

becomes comparable to specimen and/or crack size. Under the above conditions

the behavior of the material is referred to as elastic-plastic where the

simple functional relationship between the stress intensity factor and the

driving force for crack extension breaks down. The J-integral (ASTM E-813)

and crack-opening displacement (COD) (BS 5762) testing methods have been

designed on the basis of elastic-plastic fracture mechanics (EPFM) . Unfortunately, in this regime there is no geometry- independent fracture toughness

parameter, and both J and critical crack-tip opening displacement (CTOD)

,

6^, decrease with an increase in specimen size [1,2, 4, 5], Therefore, the

laboratory specimens should be designed to have plastic constraint similar to

the structural component. For example the COD testing method requires the

thickness, B, of the specimen to be equal to that of the material under

examination.

The different failure paths that generally a cracked body can follow, are

summarized in figure 1 [6], For a particular material, the change in failure

path from 1 to 5 can be associated with an increase in test temperature for a

constant specimen geometry and/or a decrease in geometric dimensions, such as

specimen thickness and crack length, and an increase in notch root radius for

a constant test temperature. The variation of fracture behavior with test

5
temperature has been traditionally characterized by ductile-to-brittle transition curves. It is well known that in body-centered cubic (bcc) metals

toughness and ductility increase with increasing temperature, and the fracture

mode changes from brittle at low temperature to ductile at high temperature.

This transition, which is a consequence of the temperature dependence of flow

stress, the temperature independence of the cleavage fracture, and the

associated notch sensitivity, was first explained by Orowan [7], In general,

there is a competition between ductile and brittle types of fracture. At a

particular temperature, a variation in parameters that change stress and

strain distributions (such as loading mode, deformation rate, and geometry)

may promote either of the fracture modes.

Most models for explaining ductile-to-brittle transitions are based on

the change in constraint at the crack tip. Recent studies [5, 6, 8, 9] have

shown that plane strain fracture toughness, K^, also reflects a ductile-tobrittle transition. The transition is associated with a change in mode of

fracture from brittle cleavage to fast ductile. Note that in the abovementioned investigations the size of the specimen has been increased with an

increase in test temperature in order to obtain ductile fracture under plane

strain and small-scale yielding conditions. For common specimen sizes most

experimental data available show that general yielding occurs before ductile

fracture initiation, and therefore one of the elastic-plastic techniques, such

as COD testing, should be applied to measure the fracture

toughness. The various regimes of fracture and deformation behavior

throughout the ductile-to-brittle transition temperature range are summarized

in figure 2. The fracture behavior can be categorized on the basis of

(a") general yielding, (b) stress state, or (c) mode of crack initiation and

propagation. Whereas (a) and (b) define the stresses and strains ahead of a

6
crack at the time of the fracture event, the fracture mode of crack initiation

and propagation, (c) , defines the fracture criterion. The shape and temperature

range of the ductile-to-brittle transition curve and the position of the

boundaries of the above-mentioned fracture behavior categories are expected to

depend on material properties (such as yield strength, work-hardening behavior,

and fracture stress), and specimen geometry.

The purpose of this investigation is to study the ductile-to-brittle

transition behavior of a shielded metal-arc weld, a typical weldment in arctic

structures. Fracture toughness was measured by the COD testing method at

temperatures throughout the ductile-to-brittle transition. The effect of

geometric dimensions on fracture toughness was studied by changing the crack

length to specimen width ratio (a/W) and the thickness of SENB specimens.

Emphasis was placed on mechanisms of crack initiation as characterized by

detailed f ractography

.

2. Materials and Experimental Procedures

2 . 1 Test Material

The base material is a 25 . 4-mm-thick plate of ABS grade EH36 steel,

composition of which is given in reference 1. A manual shielded metal-arc

welding process with 8018-C3 electrode was chosen as a typical welding practice

for structural assemblies. The welding procedure specifications are summarized

in table 1. Specimens for tensile, CVN, and COD testing were machined from

the welded plate with the length and notch orientation as shown schematically

in figure 3.

2.2 Tensile Testing

Round tensile specimens with a reduced cross section, as shown in

7
figure 4, were tested at various temperatures ranging from -258°C to 25°C.

This specimen geometry was used to enable tensile property measurement of just

the weld metal. Tensile tests were conducted in a screw-driven tensile test

machine equipped with a fixture designed for cryogenic testing at a crosshead

speed of 0.5 mm/min. Owing to the short gage length of these specimens the

strain could not be recorded directly with an extensometer . Therefore, a

calibration curve of uniform plastic strain versus time was established by

loading single tensile specimens at room temperature to various uniform strain

levels and measuring a premarked gage length of 12.7 mm in a traveling microscope with an accuracy of ±0.001 mm. This calibration curve was used to

calculate the maximum uniform strain. The total elongation was obtained by

measuring the premarked gage length after the fracture of the actual tensile

specimen.

2.3 Charpy V-Notch Testing

Standard Charpy V-notch specimens were tested in a machine with a fullscale energy value of 358 J. Testing was performed in the temperature range

-94°C to +94°C according to the ASTM-E23 standard for impact testing. Specimens were placed in a constant temperature bath (oil for temperatures above RT

and methanol for temperatures below RT) for 15 min and tested within 7 s after

removal from the bath. For each specimen the lateral expansion was measured

according to ASTM standard A-370.

2.4 COD Testing

Three-point SENB specimens of constant width (W = 21.8 mm) with various

thicknessess , B, and crack length to width ratios, a/W, (see table 2) were

prepared by fatigue precracking at room temperature according to the specifications in the British Standard for COD testing (BS 5762). The specimen geometry

and the test procedure have been explained in detail elsewhere [1] and are not

8
discussed here. The SENB specimens were tested in the temperature range

-196°C to -45°C, which resulted in critical CTOD values between 0.003 mm and

0.4 mm for establishing the ductile-to-brittle transition curves. For all the

specimens tested, macroscopically , no stable crack extension was observed on

the fracture surface. Therefore, according to the British Standard for COD

testing (B55762) the notation for the CTOD values correspond to 6 ,

designating the unstable cleavage-no prior slow crack growth situation.

However, as will be discussed later, ductile microcracking preceded brittle

fracture in some of the specimens.

As noted in the Results Section (3.1), a ductile-to-brittle transition

curve could not be obtained for specimens with a/W =0.5 because of weld

defects located in the root of the weld.

The critical CTOD values were calculated from the following relationship:

6 = 6 +6

e p

K

2

( 1—

v

2 —) + 2o E

y

r (W-a)V

E_

r (W-a)+a+Z

P

( 1 )

where

K = stress intensity factor

0^

= yield strength

E = Young's modulus

v = Poisson's ratio

r = rotational factor

P

Vp = The plastic component of the mouth-opening displacement

Z = knife edge thickness

The rotational factor, r was calculated based on the relationship

P>

between the plastic components of load-line displacement and the mouth-open i r.

displacement [10], The computed rotational factor was found to change moderately with CTOD, and the average value close to the fracture events was

g

9
approximately 0.4, the same value as suggested by the British Standard (BS 5762).

2.5 Microscopy

The fracture surfaces of the broken specimens were studied by using a

scanning electron microscope. Emphasis was placed on characterization of the

crack initiation zone in SENB specimens.

The microstructures resulting from the multiple-pass welding process and

the position of the crack tip for SENB specimens relative to these microstructures were studied by optical microscopy.

3. Results

3 . 1 Optical Microscopy

Figure 5 shows a low magnification optical micrograph of the weld. Large

globular porosities on the order of 0.5 mm in size can be seen at the root of

the weld. These defects were large enough to cause a premature fracture of

SENB specimens with a/W =0.5 geometry. The crack tip in these specimens was

located close to the porosities, and as a result, plastic deformation became

concentrated in the ligament between the crack tip and the nearest porosity.

Stresses higher than what would be expected in a porosity-free specimen were

produced in this ligament and caused a premature fracture, which was revealed

as a pop-in on the load-displacement curve. Therefore, no meaningful fracture

toughness data could be obtained for SENB specimens with a/W = 0.5 geometry.

Different weld passes can be identified in the micrograph shown in figure

5. Each inner pass contains two distinctive microstructures. One consists of

elongated prior austenite grains delineated by allotriomorph ferrite that

resulted from the solidification of the molten weld metal. The other consists

of equiaxed prior austenite grains formed by reaustenitizing of the columnar

microstructure due to the heat generated from the next weld pass. The latter

10
microstructure can be viewed as the heat affected zone of one pass due to the

following weld pass.

Figure 6 shows higher magnification micrographs of these two microstructures. The fracture toughnesses of the above-mentioned microstructures

are significantly different. As shown in figure 5, the approximate position

of the crack tip of specimens with both a/W = 0.25 and a/W = 0.75 is located

in the microstructure with equiaxed grains.

3 . 2 Tensile Properties

The tensile properties of the weld metal are given in table 3. The ultimate tensile strength and yield strength as a function of test temperature are

plotted in figure 7. The weld metal had a relatively high yield strength

(550 MPa at room temperature), as is expected for E8018-C3 electrode [11].

Both the ultimate tensile strength and yield strength increased with decreasing

test temperature. The effect of temperature on strength became more significant for test temperatures below -100°C. All the stress-strain curves showed

some indication of yield elongation but did not exhibit an upper and lower

yield point except for the specimen tested at -196°C. The increase in Luders

strain and the difference between upper and lower yield point with a decrease

in test temperature has been observed previously for bcc metals [12], The

maximum uniform strain did not change considerably with test temperature. The

work hardening rate, as reflected by the difference between yield strength and

ultimate tensile strength, decreased slightly with decreasing test temperature.

3. 3 Impact Toughness

The CVN impact energy as a function of test temperature is shown in

figure 8. The weld metal exhibited a very gradual ductile-to-brittle transition with a transition temperature range extending from -100°C to +100°('.

This wide transition temperature range was reflected in the gradual change ii

the amount of brittle, fibrous, and shear fracture observed on the f ractuiv

11
surfaces, as shown in figure 9. The Charpy V-notch data are also given in

table 4.

3 . 4 Fracture Toughness

The COD test results for the weld metal are summarized in table 5. The

critical CTOD values for the geometries studied are shown as a function of

test temperature in figure 10. As mentioned previously, due to weld defects

located at the weld root, no meaningful data were obtained for specimens with

a/W =0.5 geometry. Figure 10 shows that an increase in the specimen thickness from 12.7 mm to 25.4 mm raised the ductile-to-brittle transition temperature about 25°C. The effect of an increase in the crack length to specimen

width ratio, a/W, from 0.25 to 0.75 is revealed as a significant change in the

shape of the ductile-to-brittle transition curve towards a much steeper slope

with a slight shift to higher test temperatures. These observations are

consistent with those made for the base metal [1,2], an ABS grade EH36 steel

in normalized condition.

3.5 Fractography

3.5.1 SENB Specimens

Figures 11 and 12 show macrographs of the fracture surface for B = 25.4 mm

and B = 12.7 mm geometries, respectively. For each geometry, macrographs of

two specimens, one tested at -196°C and the other tested at a temperature

corresponding to 6^ s 0.35 mm, are presented. The important feature of these

macrographs is the relative amount of the slant or shear fracture on the sides

of the specimen. When 6^ is as large as 0.35 mm for both thicknesses it can

be seen that the amount of shear fracture is less than 10% of the total

fracture surface. Therefore, it may be concluded that the specimens tested in

this study deformed under conditions close to those of plane strain deformation and any difference in the behavior of the edges of the specimen is

12
insignificant in determining failure conditions for the specimen as a whole.

However, there has been concern whether a square fracture appearance is a

definitive indication of plane-strain conditions [9].

The fracture surfaces of the SENB specimen tested in this study did not

show any macroscopic slow crack growth preceding the unstable fracture.

However, a ductile shear zone was revealed next to the fatigue precrack,

preceding the unstable fracture, in most specimens. This shear zone should

not be confused with the stretch zone, which is commonly reported in COD

testing [13]. To clarify the various fracture zones observed on the fracture

surface, the crack-tip blunting and the fracture process at the crack tip are

discussed along with the presentation of f ractographs

.

As the load applied to a specimen containing a sharp crack is increased,

a plastic zone develops at the crack tip. The crack tip becomes blunted and

tends to be curved, as shown in figure 13b. Because of the Poisson's contraction of the material in the plastic zone, the crack has a forward movement

[13], The blunting of the crack tip is identified on the fracture surface as

a stretch zone between the fatigue crack and the area of crack extension by

cleavage or void formation. Figure 14 shows the featureless, amorphous

appearance of the stretch zone. The large geometry change at the crack tip

has been treated theoretically by modified slip-line field theory, as shown in

figure 13c. This method is justified for nonhardening materials and planestrain conditions [14]. The crack tip becomes similar to a circular notch as

a result of the progressive blunting of the crack tip with deformation.

Fracture is initiated by ductile or cleavage modes when the critical conditiofor the either of the modes is reached ahead of the crack tip. Hence, the

micromechanism of the crack initiation process can be divided into two

categories: (a) crack initiation by cleavage fracture and (b) crack initi.it i

13
by ductile fracture. Within each category the crack can extend in various

ways, as shown schematically in figures 15 and 16. In the following, crack

initiation and propagation mechanisms are discussed with respect to effects of

temperature

.

(a) Crack Initiation by Cleavage

In the lower-shelf temperature range, critical conditions for cleavage

are reached under the situation of a small plastic zone and limited blunting

at the crack tip. Assuming that cleavage is a stress-controlled fracture

process, a cleavage microcrack forms when the maximum tensile stress ahead of

* the tip reaches the fracture stress, [15,16]. It has been suggested that

the fracture stress should be reached over a characteristic distance ahead of

the crack tip [17]. This characteristic distance is related to microstructural

features that control the fracture, such as particle size and spacing and

grain size. The cleavage microcrack may join the crack tip in two ways. At _

very low temperatures, where the size of the ligament between the cleavage

microcrack and the fatigue precrack is very small and the yield strength is

high, the cleavage crack extends to the crack tip and no ductile fracture

develops (fig. 15a). This behavior was observed for SENB specimens tested at

-196°C, an example of which is shown in figure 17. Note that the stretch zone

in this case is very small 1//2 6^

= 0.002 mm) and cannot be observed on

this fractograph.

As the temperature is increased within the lower shelf energy range,

owing to a significant drop in yield strength, critical fracture stress for

cleavage is achieved at larger plastic zones accompanied by a larger stretch

zone size than in the case discussed above. In this regime, cleavage cracks

form only when a certain amount of strain hardening has occurred in the

plastic zone. Therefore, strain becomes localized in the ligament between the

14
cleavage microcrack and the fatigue crack, and the cleavage crack extends to

the notch by a shear mechanism along the maximum shear stress planes predicted

by slip-field theory [18-20]. As a result, a shear zone is formed by the

ductile (microvoid coalescence) mode, which on the fracture surface precedes

the unstable cleavage fracture (fig. 15b).

(b) Crack Initiation by Ductile Fracture

As the test temperature is raised into the ductile-to-brittle transition

temperature range, a point will eventually be reached where the combination of

stresses and strains ahead of the blunted crack favors the formation of ductile

microcracks. If the stress intensification at the tip of the ductile microcrack is large enough, a cleavage fracture will be triggered (fig. 16). In

this case the fracture surface consists of fatigue precrack, stretch zone,

shear zone, and the unstable cleavage crack, similar to the situation in

figure 15b. However, there is a definite difference between these two cases.

If the crack initiation occurs by cleavage microcracking, the shear zone size

correlates with the size of the plastic zone along the crack front, i.e. it

will be the smallest at the midthickness (plane strain) and the largest at the

edges (plane stress) . If the crack is initiated by ductile microcracking the

shear zone has a thumbnail shape with its maximum size located at the midthickness. Figures 18 and 19 show this difference in the shear zone size for

two SENB specimens. The smaller shear zone size in the midthickness in

comparison to the edges is demonstrated in figure 18. Figure 19 shows that

the shear zone attains its maximum size in the midthickness, an indication of

crack initiation by ductile fracture. Macroscopically , in the shear zone the

fracture surface is oriented 45° from the direction of the maximum tensile

stress developed by bending. Figure 20 illustrates the relative orientati"!

15
of the stretch zone and the shear zone by tilting the specimen in the SEM by

45° in two opposite directions.

If the stress intensification at the tip of the ductile microcrack is not

large enough for initiation of a cleavage crack, the crack will extend by

fibrous mode. This slow crack growth may change to an unstable fracture or it

may extend until complete fracture of the specimen.

As was mentioned in Section 3.1, the two distinct microstructures shown

in figure 6 have a significantly different toughness. This difference in

toughness was well demonstrated for one of the SENB specimens where the crack

front was located in both microstructures. The load-displacement curve for

this specimen was characterized by a pop-in that correlated with cleavage

fracture along a portion of the thickness, as marked by letter A in figure

21a. Figure 21b shows a low magnification SEM fractograph of the boundary

between the area that fractured first. A, and the area that fractured after

the arrest of the first crack, B. Note the large shear zone in the latter

area. Figure 22 shows SEM fractographs of the brittle crack in both areas.

In area B the small cleavage facets correspond to the small equiaxed grains

shown in figure 6. In area A the crack has followed the large proeutectoid

ferrite grains (a-veins; see fig. 6), as illustrated by the large columnar

cleavage facets shown in figure 22a.

3.5.2 CVN Specimens

In general, the crack initiation process in CVN specimens is very similar

to what was discussed in the preceding section [19,20] except for effects of

notch geometry. Figure 23 shows a low magnification SEM fractograph of the

CVN specimen tested at -73°C. A large shear zone next to the notch shows that

fracture was initiated by ductile fracture in this specimen. It has been shown

16
that most of the energy measured for fracture of CVN specimens is absorbed

during the process of deformation and development of this shear zone [21].

4. Comparison with Base Metal Properties

In this section, the properties of the base metal, an ABS grade EH36

steel, and the weld metal, obtained by manual metal-arc welding process with

8018-C3 electrode in accordance with table 1 specifications, are compared.

The base metal properties have been reported in references 1 and 2.

4.1 Tensile Properties

In the temperature range -196°C to 25°C, the weld metal and base metal

show a similar decrease in strength with an increase in test temperature. The

strength of the weld metal, on the the average, is about 200 MPa higher than

the base metal. The base metal has, in general, a larger total elongation

than does the weld metal. The well-defined upper and lower yield point

observed in the base metal were not exhibited by the weld metal. The absence

of yield elongation in the weld metal may be attributed to higher cooling

rates and presence of residual stresses in the weld metal. The fracture

stress for the weld metal obtained from low temperature tensile tests is 1170

MPa, which is very close to the 1156 MPa value reported for the base metal [2],

4.2 Impact Toughness

Figure 24 shows the ductile-to-brittle transition curve for both the base

metal and the weld metal. The weld metal has a much lower upper shelf energy

with a very wide transition temperature range. Note that the lower shelf

energy level is the same for both materials, consistent with the similarity in

their fracture stress values. The low resistance to ductile fracture of the

weld metal could be due to the high density of relatively fine inclusions arccarbide precipitates in the weld metal microstructure (see fig. 6). The 1-v

17
resistance to ductile fracture of the weld metal is manifested not only by the

low upper shelf energy but also by the shallow slope of the ductile-to-brittle

transition curve. For the base metal, the difference between the resistance

to brittle fracture and ductile fracture is so large that complete fracture of

the specimen occurs by either ductile or brittle fracture. The large scatter

in fracture toughness data usually observed for materials with a very narrow

ductile-to-brittle transition range is due to the statistical effects of

microstructure. For example, in the vicinity of the transition temperature of

the base metal, a specimen will break with a low absorbed energy if there is a

microstructural feature, such as carbides, positioned in a way that can

trigger cleavage fracture. On the other hand, if the crack starts in a

ductile mode, the stress level ahead of the ductile crack will never be high

enough to change to a brittle fracture mode. In comparison, a ductile crack

in the weld metal can develop under small deformation at the notch. The

degree of crack blunting at the tip of the ductile crack will be less than the

case for the base metal. Furthermore, the weld metal has a higher strength

level than the base metal does, while the fracture stress is the same.

Therefore, the ductile crack in the weld metal can change to an unstable

brittle fracture, which results in a wide range of absorbed energy and hence a

less steep ductile-to-brittle transition curve.

4.3 Fracture Toughness

The weld metal and the base metal exhibited the same KT value of 37 MPa/m

Ic

when tested at -196°C, which is consistent with the previously mentioned

observations, i.e., similar fracture stress and lower shelf impact energy. As

far as the effect of specimen thickness and crack length to specimen width

ratio on the transition curve is concerned, the weld metal and the base metal

showed the same behavior. The fracture toughness transition temperature for

18
the weld metal is about 50°C higher than for the base metal, while the slope

of the transition curve is the same in both the weld metal and the base metal

as shown in figures 25 and 26 for two different geometries. The latter

observation is quite different from the CVN impact data, which exhibited a

significant difference in slope between the weld metal and the base metal.

The inconsistency between the results of Charpy impact and fracture

toughness testing has been explained by the more complex nature of the fracture process associated with CVN testing [21]. The total CVN impact energy in

the ductile-to-brittle transition temperature range is subdivided into components associated with energy absorbed during initiation and formation of the

shear zone, energy absorbed during fibrous fracture, and energy absorbed after

the arrest of the unstable crack [20,21]. Instrumented Charpy impact testing

provides information with which the above energy components can be measured

[20,21]. The critical CTOD data are expected to correlate better with the

energy component absorbed during initiation and formation of the shear zone in

CVN specimens than with the total CVN energy. The other important factor

contributing to this inconsistency is the inhomogenity of microstructure in

the weld metal. The total impact energy is an average of the energy components absorbed during the processes of initiation and propagation of both

ductile and brittle fracture in various microstructures of the weld metal.

Furthermore, there are differences in crack geometry and loading rate between

the Charpy impact and the COD tests that contribute to the dissimilarity in

the results obtained by the two test methods.

5. Discussion

In this section, the effects of material properties and specimen yr

on the ductile-to-brittle transition temperature are discussed. Then' h

19
been many definitions for the ductile-to-brittle transition temperature. The

most common definition is the temperature, T , at which fracture occurs at

GY

the general yielding point defined as net section yielding [12], This temperature is marked by an arrow in figures 10, 25, and 26. For the limited data

available in the present investigation, this definition of the ductile-tobrittle transition temperature seems to be appropriate.

The load at general yielding for the specimen geometry used in this study

is given by:

p

gy

= l

a (W-a) 2 B 4W

( 2 )

where L is referred to as the global constraint factor. The fracture load is

calculated from the following expression:

K W

2 B (m6 a EK W

2 B

P = _c = c y (3)

F Y Y

where: Y = stress intensity coefficient = Y (a/W)

m = 2/(l-v 2

) plane strain

m = 1 plane stress

Equating P

r results in:

aJL =

E

16 m

L

2 W( 1-a/W) 4 Y

2 at T

GY

T (4)

To predict the ductile-to-brittle transition temperature for a particular

specimen geometry, assumptions have to be made about the deformation condition

20
at the crack tip (plane strain vs. plane stress), which establishes the values

of m and L, and a criterion for fracture, which establishes 6 =6 (T) . If

c c

o and E as a function of temperature are known, then the temperature T can

y

G 1

be calculated from eq (4). Unfortunately, both the deformation condition at

the crack tip and the fracture criterion may change through the ductile-tobrittle transition temperature range, which adds to the complexity of the

problem.

Based on the above discussion, the first step to be taken in explaining

the effects of material and specimen geometry on the ductile-to-brittle

transition temperature is to make assumptions about stress state and mode of

crack initiation at T . As far as the criterion for fracture is concerned,

GY

for the three SENB geometries of the weld metal tested, at T the crack has

been initiated by cleavage. This conclusion is based on detailed fractogrpahy

performed on the specimens (see section 3.5.1). It should be noted that this

conclusion is valid only for the specimen geometries tested here. For large

engineering structures it is very likely that ductile fracture can be started

at load levels much below the general yield load. The presence of less than

5% shear lips on the fracture surface of SENB specimens tested at T suggests

that crack initiation and propagation occurred under plane strain conditions

in these specimens. Such a conclusion may not be strictly correct, and there

is the possibility that square fracture can extend under a mixed mode.

Since both the weld metal and the base metal have the same microscopic

fracture stress, a^*, the 50°C difference in their transition temperature for

a given geometry (see figs. 25 and 26) should be related to the difference in

their yield strength. Assuming plane strain conditions for the deformation

near the crack tip and brittle fracture as the mode of crack initiation, it

can be shown that [22]:

21
where ct„ and a are the nominal stresses at the fracture and general yielding

point, respectively, and g is a parameter that accounts for finite specimen

and crack geometry. Equation (5) has been derived by equating the plastic

zone size of the Dugdale model to the plastic zone size of the slip-field

theory at the point of the fracture. Equation (5) implies that for a given

specimen geometry equivalent values of o / a lead to equivalent values of

r GY

;

Oj. /o . Therefore, for a constant o^/o^ ratio, as arbitrarily chosen for the

I y r GY

definition of the transition temperature:

or

constant

Of (material 1) a

^ (material 2)

a o y y

( 6 )

(7)

For the base metal and the weld metal:

a

* (1) = a

f * (2) = 1160 MPa

Therefore

a (1) = a (2) (8)

y y

For the B = 25.4 mm and a/W = 0.25 specimen geometry of the weld metal,

(T = -83°C) = 590 MPa. The temperature at which the yield strength of the

GY

22
base metal is equal to 590 MPa is -132°C [2]. For the B = 12.7 mm and a/W =

0.25 geometry for the weld metal a = 640 MPa at T^ = -108°C, which leads to

a T of -150°C for the base metal. The transition temperature shifts of 49°C

GY

and 42°C predicted for B = 25.4 mm and B = 12.7 mm specimen geometries correlate

surprisingly well with the experimental results.

For SENB specimens with load span of 4W, the parameter g in eq (5) is

given by the following expression [22]:

g

1 2 TT 3

W( 1-a/W)

Ti p y

(9)

where p is the notch root radius. In the case of sharp cracks, p becomes P q

,

which is a microstructural characteristic distance [24]. For a constant

A- 2 specimen width, an increase in the a/W ratio decreases (1-a/W) y , which leads

to a larger o , according to eq (5). Consequently, a decrease in transition

temperature with an increase in a/W is predicted, which is contradictory to

the experimental results. However, eq (5) was derived for a constant constraint

4 2 at the crack tip. Note that the term (1-a/W) y also appears in eq (4). It

seems the only way to explain the increase in transition temperature due to an

increase in a/W ratio is through the constraint model. Specimens with larger

a/W ratios produce higher plastic constraints, i.e., larger L values, and

4 2 therefore compensate for the change in the (1-a/W) y term. The steep slope

of the transition curve for a/W = 0.75 geometry has been attributed to the

relaxation of the crack-tip constraint as the plastic zone approaches a free

surface [1,2],

The effect of thickness on the ductile-to-brittle transition temperature

has been explained either by constraint effects [1,2, 4, 5] or by statistical

23
effects [23]. The former model attributes the decrease in transition temperature due to a reduction of specimen thickness to a loss in stress triaxiality

at the crack tip, which reduces the effective flow stress in the plastic zone.

On the basis of the statistical model, the effect of thickness is explained by

the probability of the formation of crack nuclei within the volume of material

sampled along the crack front. Therefore, the statistical model predicts a

greater scatter in fracture toughness data for thin specimens than for thick

specimens and a decrease in mean fracture toughness at a given temperature

with increasing specimen thickness. The larger scatter in toughness data as a

function of test temperature observed for the 12. 7-mm-thick specimens (see

figs. 25 and 26) may demonstrate the validity of the statistical model.

However, with the limited data available, the effect of specimen thickness on

the transition temperature cannot be strictly attributed to explanations of

either of the above mentioned models. Most probably, both the constraint

effects and statistical effects contribute to the shift in the transition

temperature.

So far, the effects of thickness, material properties, and crack depth on

the ductile-to-brittle transition temperature have been discussed. As far as

design problems are concerned, specimen thickness seems to be the most important

of these three variables; however, no well established theoretical treatment

of its effect is known. Both constraint and statistical models predict a

saturation thickness for a particular material, so that a further increase in

specimen thickness will have no effect on measured fracture toughness. The

results of this investigation indicate that at temperatures above T^y fracture

is initiated by ductile fracture. This observation implies that the critical

CTOD in the transition range is controlled by the subsequent propagation of a

24
brittle crack from a ductile microcrack rather than the initial fatigue crack.

Therefore, simplified theoretical treatments similar to what was given here,

although qualitatively correct, may not be adequate for design purposes. It

is suggested that the mechanisms that control the ductile-to-brittle transition

curves are very complicated and a better understanding of stress and strain

distributions and micromechanism of the fracture process ahead of a progressive

blunting crack is required.

6. Summary and Conclusions

The ductile-to-brittle transition behavior of a shielded metal-arc weld,

a typical steel weldment in arctic structures, has been studied and compared

to that of the base metal, an ABS grade EH 36 steel in normalized condition.

The toughness and the ductile-to-brittle transition temperature of the weld

metal and the base metal measured by different test methods for various

specimen geometries are summarized in table 6. On the basis of the results

obtained from tensile, Charpy V-notch, and fracture toughness testing and also

scanning and optical microscopy, the following conclusions have been reached:

1. Specimen dimensions have a significant effect on the ductile-tobrittle transition curve. An increase in the specimen thickness

from 12.7 mm to 25.4 mm for a/W =0.2 geometry resulted in an upward

shift of 25°C in the ductile-to-brittle transition temperature.

These results are consistent with the observations made for the base

metal [1,2].

2. In the transition region, crack initiation occurs by ductile fracture

even for those specimens that do not show macroscopic slow crack

25
growth. The ductile crack initiation was detected as a shear zone

preceding the cleavage crack on the fracture surface.

3. The weld metal and the base metal have equal resistance to cleavage

fracture, as indicated by equal K.^, tensile fracture stress, and

CVN lower shelf energy values measured at low temperatures. However,

the weld metal has a higher fracture toughness transition temperature, a result associated with the high strength and low resistance

to ductile fracture of the weld metal.

4. The slope of the fracture toughness transition curve is the same in

both the weld metal and the base metal. However, the weld metal has

a much wider CVN transition temperature range than the base metal.

The inconsistency between the results of CVN impact and fracture

toughness testing in the ductile-to-brittle transition temperature

range is suggested to be associated with the more complex fracture

induced by CVN testing and the inhomogeneity of the weld metal

microstructure.

5. There are two distinctly different microstructures in each weld

pass. One consists of columnar prior austenite and acicular ferrite

grain, and the other consists of equiaxed ferrite grains formed by

reaustenitizing of the columnar microstructure. The results of this

study suggest that the former microstructure has a lower fracture

toughness and, when possible, the cleavage crack follows the blocky

ferrite (a-veins) formed at the prior austenite grain boundaries.

26
7. References

[1] T. L. Anderson and H. I. McHenry: "Fracture Toughness of Steel Weldments

for Arctic Structures," NBSIR83-1680, National Bureau of Standards,

Boulder, Colorado, 1982.

[2] T. L. Anderson: "The Effect of Crack-tip Region Constraint on Fracture

in the Ductile-to-Brittle Transition Region", Ph.D. Thesis //T-2741,

Colorado School of Mines, Golden, Colorado, 1983.

[3] J. F. Knott: "Fundamentals of Fracture Mechanics", John Wiley and Sons,

Inc., 1973.

[ 4 ] G. Green and J. F. Knott: J. Mech. Phys. Solids., 1975, vol. 23, p. 167.

[5] M. G. Dawes: in "Elastic-plastic Fracture", ASTM STP 668, J. D. Landes,

J. A. Begley, and G. A. Clarke, Eds., American Society for Testing and

Materials, 1979, p. 307.

[6] I. Milne and G. G. Chell: in "Elastic-plastic Fracture", ASTM STP668

J. D. Landes, J. A. Begley, and G. A. Clarke, Eds., American Society for

Testing and Materials, 1979, p. 358.

[7] G. E. Orowan: Rep. Prog. Phys., 1948, vol. 12, p. 185.

[8] S. T. Rolfe and J. M. Barsom: "Fracture and Fatigue Control in

Structures", Prentice-Hall, 1977, p. 92.

[9] B. Marander and G. Sanz: in ASTM STP 631, 1976, p. 72.

[lO ] I.-H. Lin, T. L. Anderson, R. deWit and M. G. Dawes: Int. J. Fracture,

1982, vol. 20, p. R3.

[11] Metals Handbook, "Welding and Brazing", eighth edition, vol. 6, p. 191.

[12] D. McLean: "Mechanical Properties of Metals", John Wiley and Sons, 1962,

p. 209.

[13] J* D. Harrison: "The State-of-the-Art in Crack Tip Opening Displacement

(CTOD) Testing and Analysis", Report #7302.10/80/210.2, The Welding

Institute, Abington, Cambridgeshire, England.

[14] J* R* Rice and M. A. Johnson: in "Inelastic Behavior of Solids", M. F. Kannien, et al., Eds., McGraw Hill, 1970, p. 641.

[15] T. R. Wilshaw, C. A. Rau, and A. S. Tetelman: Eng. Fract. Mech., 1968,

vol. 1, p. 191.

[l6 ] 3* Malkin and A. S. Tetelman: Eng. Fract. Mech., 1971, vol. 3, p. 151.

27
[17] R. 0. Ritchie, J. F. Knott, and J. R. Rice: J. Mech. Phys . Solids, 1973,

vol. 21, p. 395.

[18] J. F. Knott: Metal Sci. , 1980, vol. 14, p. 327.

[19] F. Zia-Ebrahimi, D. K. Matlock, and G. Krauss: Scripta Metall., 1982,

vol. 16, p. 987.

[20] F. Zia-Ebrahimi and G. Krauss: Acta Metall., in press.

[21] F. Zia-Ebrahimi and G. Krauss: Met. Trans. A., 1983, vol. 14, p. 1109.

[22] A. S. Tetelman and A. J. McEvily, Jr.: "Fracture of Structural Materials",

John Wiley and Sons, 1967, p. 313.

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

[24] R. 0. Ritchie, B. Francis and W. L. Server: Met. Trans. A., 1976, vol.

7, p. 831.

28
Table 1. Specifications for welding procedure

Type of process

Preheat

Electrode class and

size

Welding position

Voltage, current (DCRP)

and velocity

Manual shielded metal arc

None

E80318, 4.8, 4.0, and 3.2

Flat

4.8 mm: 28 V, 210 A

100-150 mm/min

4.0 mm: 25V, 170 A

100-125 mm/min

3.2 mm: 23 V, 120 A

75 mm/min

Joint design and

bead placement

mm

29
Table 2. Test matrix for fracture toughness tests of welded SENB specimens

a/W B = 12.7 mm B = 25.4 mm

0.25 X X

0.5 X

0.75 X

W = 21.8 mm

30
Test

Yield

Ultimate

Uniform

Total

Fracture

Reduction

Temperature

Strength

Tensile

Strength

Strain

Strain*

Stress

in

Area

(°C)

(MPa)

(MPa)

(%)

(%)

(MPa)

(%)

l in cni co .—i r^. n- <r i I

I n- n- r-~ cn co | |

co vO r^.

1 1 r^ 00 m o CN VO m vO CN <r CN i—

«—1 i—

i—H i—rH i—l *—

r—1 r

1 H

ocncNOcNor-'in l I CO CO CO CO CO <r —' CN | i

00 LO o On o

t—

T—

CN CN CO 00 CN

*—

I—

1—1 i—

i—

r—

Ln o o t—

o nCO 00 cn CN O o CN

v£> 00 O CN i

1 H

o o 00 CO

CN r-~ o m cn cn Ln

Ln Ln m id 00

CNpH

II

Ln «-H co oo co

CN CO CO »—i ^

*

*K toe

CO vO *—

oo (U a> <r in t—

i—

*—

CN CN

1 l i 1

*Gage

31

'''Premature

fracture

due

to

the

existence

of

welding

defects

in

the

gage

length
Table 4. Charpy V-notch data for the weld metal

Test Temperature

(°C)

Absorbed Energy

(J)

Lateral Expansion

(mm)

94 203 2.137

58 182 2.180

25 161 1.948

1 126 1.644

-6 122 1.619

-24 80 1.145

-35 89 1.218

-45 80 1.123

-73 27 0.424

-94 8.8 0.185

32
Table

5.

Critical

CTOD

data

for

the

weld

metal

m 00m

UB 00 *<r cn

to

g CN CN o o '— • • • • o o o o

CO t—

40 00 c 3 m 00 »—

aj

g

g ng g CO • • • • •H o o o o u <r oo

QJ . . a. lo —

cn CM CN

ii ii ii /-N * O 00 o cn

m

w « s H o 40

m 40 i 1 i 1

04 o 04 04 00 1—CN cn

OB

m ao o 04 40 CN o

°o

g cn CN CN| »“ i o o o o • • o o o o o o o o o

CAC

qj e g egg •H o cn 40 cn cn CO CN o <r oo 3 40 04 ON 04 00 CO

Q) . . CN CN CN CN CNI CNI CN CN CM

Q* in • cd CO CNI CNI o o o o o O o o o

II II II

ca « S +

CN 40 00 CN cn 40 40

U

m 40 40 40 40 00 o 04

H o i 1 1 1 1 i 1 r—

. 1 1 1

40 04 cn i—

o 04 00 cn

OB OO 40 CN 40 *—1 cn o to g co CNI cn *—

CN o o cnc o o o o o o o o QJ g S BBS •Hu r-. oo

aj . • a- cn ^ 40 i—

1—

i—

00 <}• CN

on i—i cm 3 *<r <d" cn cn cn cn — CNJ CN CN CN CN CN CN CN H H n cO • o o o o o o o o C CQ S K /—\

-K +

CJ oo 00 cn 40 r- CN 00 r—

H o 40 40 00 00 04 o ON 1 1 i 1 1 1 1 1 eo

•H

4-1 CO3

tO

3.X

QJ

l-i O

4-1X

cu

qj 4=

QJcu

cn

0)

S-t 3

j—

u33

COo3u

•HBC

co

QJOc

QJM

QJ

4-1 •HXI

QJ X

QJ3

X)C

•HIa.o

D.c

t—

K

33

**Five

times

higher

displacement

rate

(4

mm/min)
Table 6. Summary of the toughness and the ductile-to-brittle transition

temperature data for the base metal and the weld metal

Conditions Parameter Base Metal Weld Metal

T =-196°C Kt (MPa /m) ic

37 37

T = 60°C CVN Energy (J) 335 47.8

CVN Energy = 27 J T (°C) -78 -74

B = 12.7 mm t

gy

<° c > k

-140 -98

a/W = 0.2-0.25 T (°C) -128 -79

B = 25.4 mm t

gy

( °S>

-130 -83

a/W = 0.75 T (°C) -111 -59

B = 25.4 mm t

gy & -120 -55

a/W = 0.2-0.25 T (°C) -100 -50

6 = 0 . 2 mm

c

34
Figure 1 The various failure path that a cracked body can follow.

35
aoio ivoiiido

coai>

5-1 3

O

Co

•H

COc

cd

5-1

<U

CO

ai

5-1 Pcu

5-1 O

cde

cuu

CO

CN

a)

5

-( 3M

•H

Pt-I

36

TEST

TEMPERATURE
Rolling Direction

Figure 3

a) SENS Specimen

b) CVN Specimen

c) Tensile Specimen

The orientation of the test specimens relative to the

steel weldment.

37
38

Figure

4.

Dimensions

of

the

tensile

specimens.
39

i>;ure

5.

Low

magnification

optical

micrograph

of

the

weldment.

The

approximate

position

of

the

crack

tip

for

a/W

values

are

shown.
Figure 6. Two distinctively different microstructures of the weld

metal. (a) Columnar prior austentite grains, blocky ferrite

at the grain boundaries and acicular ferrite. (b) Fine

equixed ferrite grains.

40
1200

BdW 'SS3d±S 0NIU33NI0N3

0)Vi3

4-1

Cti Vi

<Ucv£

U)c

<u

T3c

03 "O

(U

V

3

41

TEST

TEMPERATURE
qi-U 'ADH3N3 ±OVdWI NAO

r ‘ADU3N3 lOVdWI NAO

QJ

S-i 3

4-

>

5-

i 1)P-e

cu

2

>

u

00

a)

S-i P

M)

•H

f*-i

42

TEST

TEMPERATURE,

°C
43

Figure

9.

Macrographs

of

the

fracture

surfaces

of

CVN

specimens

tested

in

the

transition

temperature

range.
^dodo CO CNI T- o CM

1

uiuj *aoio nvouido

CO

(1) 3

3

4-1 3

3

CD

a) a

a) e 3 CD ,3 P

4-1 3

3 O

O •H

<4-4 U

•H

U bO

'—I 3

cfl a) o g u a

•h <u

s-i a* u cn

a)3

3

hO •H

44

TEST

TEMPERATURE
45

Figure

11.

Macrograph

of

the

fracture

surfaces

of

SENB

specimens

B=25.4

geometry

tested

as

(a)

-196°C

and

(b)

-52°C.
46

Figure

12.

Macrographs

of

the

fracture

surface

of

SENB

specimens

with

B=12.7

mm

geometry

tested

at

(a)

-196°C

and

(b)

-68
a)

Sharp

crack 4 7

(a)

Stationary

sharp

crack,

(b)

blunted

crack

under

the

applied

load,

formation

of

the

stretch

zone,

and

(c)

development

of

the

slip-filed

lines

due

to

the

large

geometry

change

at

the

crack

tip.
48

Figure

14.

SEM

fractograph

showing

the

fatigue

crack,

stretch

zone,

shear

zone

and

the

unstable

brittle

crack

for

a

SENB

specimen

with

B=25.4

mm

geometry

tested

at

-74°C,
Unstable

brittle

49

1

igure

15.

Schematic

showing

crack

extenstion

by

cleavage

crack

initiation

mechanism.
50

Figure

16

.

Schematic

showing

crack

extension

by

ductile

crack

initiation

mechanism.
51

Figure

17.

SEM

fractograph

showing

the

absence

of

the

shear

zone

at

the

crack

tip

for

SENB

specimen

with

B=25.4

mm

geometry

tested

at

-196°C.
Shear Zone

Figure 18. SEM fractographs showing the difference in the size of the

sheat zone formed at (a) the midthickness and (b) the

edge of the specimen. SENB specimen with B=25.4 mm geometry

tested at -74°C.

52
Figure 19. SEM fractographs showing the difference in the size of the

shear zone formed at (a) the midthickness and (b) the

edge of the specimen. SENB specimen with B=12.7 mm geometr

tested at -68°C.

53
f

Unstable

Brittle Crack

Shear Zone

Stretch Zone

f

Fatigue Crack

1

Unstable

Brittle Crack

Shear Zone

Fatigue Crack

50 yum

Figure 20. SEM fractographs showing the relative orientation of the

stretch zone and the shear zone, (a) 0° tilt

(b) -45° tilt.

54
r

Unstable Brittle

Crack (UBC)

Stretch Zone

Fatigue Crack

<FC)

L

Figure 20. (continued)

(c) +45° tilt (d) A schematic showing the specimen at 4 5°

tilt angle relative to the electron beam.

55
Figure 21. (a) Macrograph of the fracture surface of the SENB specimen with

a/W = 0.75 geometry tested at -48°C. (b) Low magnef ication SEM

fractorgraph showing the boundary between areas A and B as marked on the macrograph shown in (a)

.

56
Figure 22. SEM fractograph showing the brittle fracture in areas A (a) am!

B (b) as marked on the macrograph shown in figure 21.

57
58

Figure

23.

SEM

fractograph

showing

shear

zone

next

to

the

notch

of

CVN

specimen

tested

at

-73°C.
r

‘ A9H3N3 lOVdm NAD

59

TEST

TEMPERATURE
o

0

l

o

1

o

0

T“

1

0

in

1

o

0

CsJ

1

O

ujuj 4 aoio nvoiimo

S-l 4-1 <UBo

qj00c

a)B

•HO

qj

a3 03

S-4 4J

4-J OJB a

O T3

EH i—

O QJ5

03 QJ a

rl 4-1 +J •H 4-1 H O U

LO

CXI0) H

3

00

"H

60

TEST

TEMPERATURE
wiu ‘aoio ivouido

a) x.

•u

4-1Ou

•u

a)Bo

0) 00

c Q CO O

E—1 i—

O CO u

rH <U

co e o

•H TJ

±J r—

(

•H 0) M 3 O

vO

CN

a)3

00

u.

61

TEST

TEMPERATURE
NBS-114A (REV. 2-80

U.s. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO. NBSIR 85-3020

2. Performing Organ. Report No, 3. Publication Date

April 1985

4. TITLE AND SUBTITLE

Ductile-to-Brittle Transition in Steel Weldments for Arctic Structures

5. AUTHOR(S)

F. Zia-Ebrahimi

6. PERFORMING ORGANIZATION (If joint or other than N BS. see in struction s) 7. Contract/Grant No. national bureau of standards

DEPARTMENT OF COMMERCE 8. Type of Report & Period Covered

WASHINGTON, D.C. 20234

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

U.S. Department of Interior

Minerals Management Service

12203 Sunrise Valley Drive

Reston, Virginia 22091

10. SUPPLEMENTARY NOTES

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

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

bibliography or literature s urvey^mention it here) This report s umma rizes the work performed in support of the development of fracture criteria for steel weldments in arctic structures. The ductile-to=brittle transition behavior of a shielded metal-arc weld, typical of steel weldments in arctic structures, has been studied. Fracture toughness, Charpy V-notch impact energy, and tensile

properties have been measured as a function of temperature throughout the ductile-tobrittle transition range. The effect of geometric dimensions on fracture toughness has

been studied for three geometries of single-edge-notch-bend (SENB) specimens. The

fracture surfaces of broken specimens have been characterized by scanning electron

microscopy (SEM) . The microstructure of the multiple-pass weldment has been studied

by optical microscopy. The mechanical properties of the steel weldment have been

compared to the base metal, an ABS grade EH36 steel in normalized condition.

The results indicate that the resistance to brittle fracture is the same in both

the base metal and the weld metal; however, the weld metal has a much lower resistance

to ductile fracture than the base metal does. The high ductile-to-brittle transition

temperature of the weld metal is attributed to its high tensile strength and low re- sistance to ductile fracture.

The ductile-to-brittle transition temperature increased by 25°C with increasing

specimen thickness (from 12.7 mm to 25.4 mm), an effect that can be attributed to both

the constraint and statistical effects. The slight increase in ductile-to-brittle

transition temperature with increasing crack length/specimen width ratio is discussed

on the basis of constraint effects.

The SEM fractography revealed a shear zone preceding the unstable brittle fracturel in SENB specimens tested in the transition temperature regime. This finding shows that! tne fracture initiation process in these specimens is controlled by ductile fracture

rather than brittle fracture, although no macroscopic slow crack growth was observe?.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons)

crack initiation; fracture toughness; fracture surface; microstructure; 3-point band

tests; transition temperature; weldments

13. AVAILABILITY

[X] Uni imited

| | For Official Distribution. Do Not Release to NTIS

[H Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402.

QjQ Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

68

15. Price

USCOMM-DC 6043-P80

* U.S. GOVERNMENT PRINTING OFFICE: 1985-578-736/25191 REGION NO. 8