Seminar on Brittle and Ductile Fracture

Seminar on Brittle and

Ductile Fracture

Department of Mechanical Engineering

Manipal Institute of Technology

2008

Chetan Purushottam Bhat

Mtech (CAMDA)

Seminar on Brittle and Ductile Fracture



Contents

Introduction

Mechanism of Ductile Fracture

Mechanism of Brittle Transgranular Fracture (Cleavage)

Intergranular Fracture

Ductile to Brittle transition

Notched-bar Impact Tests

Ductile to Brittle Transition-Temperature Curve (DBTT)

Criterion for Transition Temperature

Metallurgical Factors affecting Transition Temperature

Conclusion

References




Introduction

Since the World War II there has been great progress in understanding the ways in which

the materials fracture. Nevertheless, it is still not possible to use this knowledge, together

with other material properties, for predicting fracture behaviour in engineering terms with a

high degree of confidence. Metallic materials, especially alloys are highly complex. An

indication of this complexity is given by the figure below, which shows various

microstructural features (not all of which need to be present in a material) and also the two

main types of fracture path, transgranular and intergranular. Of fundamental importance is

the fact that almost all the structural materials are polycrystalline, i.e. they consist of

aggregate of grains, each of which has a particular crystal orientation. The only exceptions

are single crystal turbine blades for high performance jet engines.

.

Schematic of microstructural features in metallic materials. Ref [1]

Metals fail by two broad classes of mechanisms: Brittle and Ductile failure

The Brittle fracture has following characteristics:

• There is no gross plastic deformation of the material and failure occurs with low

energy absorption.

• The surface of the brittle fracture tends to be perpendicular to the principal tensile

stress although other components of stress can be factors.

• Characteristic crack advance markings frequently point to where the fracture

originated.

• The path the crack follows depends on the material's structure. In metals,

transgranular and intergranular cleavage are important. For energy related reasons,

a crack will tend to take the path of least resistance.




The Ductile fracture has the following characteristics

• There is considerable deformation before failure and lot of energy is required

compared to ductile failure.

• The fracture surface is dull and fibrous. The appearance of a ductile fracture at a high

magnification is a surface with indentation as if marked by an ice-cream scooper.

This surface morphology is appropriately called dimpled

• Rupture by total necking is very rare because most metals contain second phase

particles that act as initiation sites for void. However high purity metals such as

copper nickel gold and other very ductile material fail with very high reduction in

areas.

• Most structural material exhibit considerable strain before reaching the tensile or

ultimate strength.

Schematic classification of fracture processes. Ref [2]

Comparison of stress strain curve of pure ductile moderately ductile brittle

brittle and ductile material fracture fracture fracture




Mechanism of Ductile Fracture

Ductile fracture is caused by overload and depending on constraint can often be recognised

immediately from macroscopic examination of failed specimen or component .If there is

very little constraint there will be a significant amount of contraction before failure occurs.

When there is high constraint (e.g. thick sections) a ductile fracture may occur without

noticeable contraction. In such cases the only macroscopic difference is the reflectivity of

the fracture surface, which tends to be dull for a ductile fracture and shiny and faceted for a

brittle fracture.

The figure below schematically illustrates the uniaxial tensile behaviour in a ductile metal.

The material eventually reaches an instability point, where strain hardening cannot keep

pace with loss in cross sectional area, and a necked region forms beyond the maximum load.

In very high purity materials, the tensile specimen may neck down to a sharp point, resulting

in extremely large local plastic strains and nearly 100% reduction in area. Materials that

contain impurities, however, fail at much lower strains. Microvoids nucleate at inclusions

and second phase particles; the voids grow together to form a macroscopic flaw, which

leads to fracture.

Uniaxial tensile deformation of ductile material. Ref [3]

In ductile materials, the role of plastic deformation is very important. The important feature

is the flexibility of slip. Dislocations can move on a large number of slip systems and even

cross from one plane to another (in cross-slip). Consider the deformation of a single crystal

of copper, a ductile metal, under uniaxial tension. The single crystal undergoes slip

throughout its section. There is no nucleation of cracks, and the crystal deforms plastically

until the start of plastic instability, called necking. From this point onward, the deformation

is concentrated in the region of plastic instability until the crystal separates along a line or a

point. In the case of a cylindrical sample, a soft single crystal of a metal such as copper will

reduce to a point fracture. Figure below shows an example of such a fracture in a single

crystal of copper.




A point fracture in soft single crystal sample of copper. Ref [2]

In crystalline solids, cracks can be nucleated by the grouping of dislocations piled up against

a barrier; such cracks are called Zener-Stroh cracks. High stresses at the head of a pileup are

relaxed by crack nucleation, as shown in Figure. But this would occur only in the case where

there is no relaxation of stresses by the movement of dislocations on the other side of the

barrier.

Grouping of dislocations piled up at a barrier and leading to the formation of microcracks. Ref [2]

The most familiar example of ductile fracture is that in uniaxial tension, giving the classic

"cup and cone" fracture. An ideal plastic material in which no strain hardening occurs would

become unstable in tension and begin to neck just as soon as yielding took place. However,

a real metal undergoes strain hardening, which tends to increase the load carrying capacity

of the specimen as deformation increases. This effect is opposed by the gradual decrease in

cross sectional area of the specimen as it elongates. Necking or localized deformation begins

at maximum load, where there is increase in stress due to decrease in the cross sectional

area of the specimen and becomes greater than the load carrying ability of the metal due

strain hardening. When the maximum load is reached, the plastic deformation in a

cylindrical tensile test piece becomes macroscopically heterogeneous and is concentrated in

a small region. This phenomenon is called necking .The final fracture occurs in this necked

region and has the characteristic appearance of a conical region on the periphery resulting

from shear and a central flat region resulting from the voids created there.

In practice, materials generally contain a large quantity of dispersed phases. These can be

very small particles (1 to 20 nm) such as carbides of alloy elements, particles of intermediate

size (50 to 500 nm) such as alloy element compounds (carbides, nitrides, carbonitrides) in

steels, or dispersions such as AI2O3, in aluminium and ThO2, in nickel. Precipitate particles

obtained by appropriate heat treatment also form part -of this class (eg., an Al-cu-Mg




system), as do inclusions of large size (on the order of millimetres)-for example, oxides and

sulphides. If the second-phase particles are brittle and the matrix is ductile, the former will

not be able to accommodate the large plastic strains of the matrix, and consequently, these

brittle particles will break in the very beginning of plastic deformation. In case the particle &

matrix interface is very weak, interfacial separation will occur. In both cases microcavities

are nucleated at these sites. Generally, the voids nucleate after a few percent of plastic

deformation, while the final separation may occur at around 25%.The microcavities grow

with slip, and the material between the cavities can be visualized as a small tensile test

piece. The material between the voids undergoes necking on a microscopic scale, and the

voids join together. However, these microscopic necks do not contribute significantly to the

total elongation of the material. This mechanism of initiation, growth, and coalescence of

microcavities gives the fracture surface a characteristic appearance. When viewed in the

scanning electron microscope, such a fracture appears to consist of small dimples, which

represent the microcavities after coalescence. In many of these dimples, one can see the

inclusions that were responsible for the void nucleation.

Dimple shape is strongly influenced by the type of loading. This is illustrated in figure.

Fracture under local uniaxial tensile loading usually results in formation of equiaxed

dimples. Failures caused by shear will produce elongated or parabolic shaped dimples that

point in opposite directions on matching fracture surfaces. And tensile tearing produces

elongated dimples that point in the same direction on matching fracture surfaces.

Dimple formation owing to uniaxial tensile loading, shear and tensile tearing. Ref [1]




Fracture by Void Nucleation, Growth, and Coalescence.

We describe the process of fracture by void nucleation, growth, and coalescence in some

detail because of its great importance in metals. In materials where the second phase

particles and inclusions are well bonded to the matrix, void nucleation is often the critical

step, fracture occur soon after the voids form. When void nucleation occurs with little

difficulty, the fracture properties are controlled by the growth and coalescence of voids, the

growing voids reaches critical size, relative to their spacing and a local plastic instability

develops between voids, resulting in failure.

Void Nucleation

A void forms around a second phase particle or inclusion when sufficient stress is applied to

break the interfacial bonds between the particle and the matrix. A number of models for

estimating void nucleation stress have been published, some of which are based on

continuum theory, while others incorporate dislocation-particle interactions. The latter

models are required for particles < I µm in diameter.

The most widely used continuum model for void nucleation is due to Argon, et al. They

argued that the interfacial stress at a cylindrical particle is approximately equal to the sum

of the mean (hydrostatic) stress and the effective (von Mises) stress. The de-cohesion stress

is defined as a critical combination of these two stresses:

And , and are the principal normal stresses. According to the Argon, et al model, the

nucleation strain decreases as the hydrostatic stress increases. That is, void nucleation

occurs more readily in a triaxial tensile stress field, a result that is consistent with

experimental observations.

Void Growth and Coalescence

Once voids form, further plastic strain and hydrostatic stress cause the voids to grow and

eventually coalesce. If the initial volume fraction of voids is low (< l0%),each void can be

assumed to grow independently; upon further growth, neighbouring voids interact. Plastic

strain is concentrated along a sheet of voids, and local necking instabilities develop. The

orientation of the fracture path depends on the stress state

Many materials contain a bimodal or trimodal distribution of particles. For example, a

precipitation-hardened aluminium alloy may contain relatively large intermetallic particles,

together with a fine dispersion of submicron second phase precipitates. These alloys also




contain micron-size dispersoid particles for grain refinement. Voids form much more readily

in the inclusions, but the smaller particles can contribute in certain cases. Bimodal particle

distributions can lead to so-called "shear" fracture surfaces, as described below. Figure

illustrates the formation of the "cup and cone" fracture surface that is commonly observed

in uniaxial tensile tests. The neck produces a triaxial stress state in the centre of the

specimen, which promotes void nucleation and growth in the larger particles. Upon further

strain, the voids coalesce, resulting in a penny-shaped flaw. The outer ring of the specimen

contains relatively few voids, because the hydrostatic stress is lower than in the centre. The

penny-shaped flaw produces deformation bands at 45o from the tensile axis. This

concentration of strain provides sufficient plasticity to nucleate voids in the smaller more

numerous particles. Since the small particles are closely spaced, instability occurs soon after

these smaller voids form, resulting in total fracture of the specimen and the cup and cone

appearance of the matching surfaces. The central region of the fracture surface has a

fibrous appearance at low magnifications, but the outer region is relatively smooth. Because

the latter surface is oriented 45o from the tensile axis and there is little evidence (at low

magnifications) of microvoid coalescence, many refer to this type of surface as shear

fracture. The 45o angle between the fracture plane and the applied stress results in a

combined Mode I/Mode II loading.







Mechanism of Brittle Transgranular Fracture (Cleavage)

Figure shows an interesting example of cleavage fracture in chromium hard plating on a steel shaft.

A truly brittle fracture is caused by cleavage. The term brittle fracture can be misleading.

Since essentially ductile fracture (microvoid coalescence) under high constraint may show

the same lack of contraction expected for cleavage. Cleavage generally takes place by the

separation of atomic bonds along well-defined crystal planes. Ideally, a cleavage fracture

would have perfectly matching faces and be completely flat and featureless. However,

structural materials are characteristically, polycrystalline with the grains more or less

randomly oriented with respect to each other. Thus cleavage propagating through one grain

will probably have to change direction as it crosses a grain or sub-grain boundary (sub-grains

are regions within a grain that differ slightly in crystal orientation). Such changes in

direction, results in the faceted fracture surface. In addition most structural material contain

particles, precipitates or other imperfections that further complicate the fracture path, so

that truly featureless cleavage is rare, even within a single grain or subgrain. The changes of

orientation between grains and subgrain’s and the various imperfections produce markings

on the fracture surface that are characteristically associated with cleavage.

Figure illustrates some typical feature associated with cleavage. A principal feature is river

pattern, which are steps between cleavages on parallel planes. Rivers patterns always

converge in the direction of local crack propagation. If the grains or subgrains are connected

by a tilt boundary, which means that they are misoriented about a common axis, the river

patterns are continuous about the boundary. But if adjacent grains or subgrain’s are axially

misoriented i.e. they are connected by a twist boundary, the river patterns don’t cross the




boundary but originate at it. Besides river patterns distinct feature of cleavage is feather

markings. The apex of these fan-like markings point back to the fracture origin, and

therefore this feature can also be used to determine the local direction of crack

propagation.

Typical features associated with cleavage

We mentioned that cleavage occurs along specific crystallographic planes. As in a

polycrystalline material, the adjacent grains have different orientations; the cleavage crack

changes direction at the grain boundary in order to continue along the given

crystallographic planes The cleavage facets seen through the grains have a high reflectivity,

which gives the fracture surface a shiny appearance. Sometimes the cleavage fracture

surface shows some small irregularities-for example, the river markings. What happens is

that, within a grain, cracks may grow simultaneously on two parallel crystallographic planes

the two parallel cracks can then join together, by secondary cleavage or by shear, to form a

step. Cleavage steps can be initiated by the passage of a screw dislocation. In general, the

cleavage step will be parallel to the crack's direction of propagation and perpendicular to

the plane containing the crack. As this configuration, would minimize the energy for the step

formation by creating a minimum of additional surface. A large number of cleavage steps

can join and form a multiple step. On the other hand, steps of opposite signs can join and

disappearing. The junction of cleavage steps results in a figure of a river and its tributaries.

River markings can appear by the passage of a grain boundary as shown in. We know that

cleavage crack tends to propagate along a specific crystallographic plane. This being so,




Many small regions of river patterns are apparent in Figure

Figure shows nice river patterns (twist misorientation) at a higher magnification and also shows tilt

boundaries, where the grains are merely tilted with respect to each other.




when a crack passes through a grain boundary, it has to propagate in a grain with a different

orientation. Figure shows the encounter of a cleavage crack with a grain boundary. After

they meet, the crack should propagate on a cleavage plane that is oriented in a different

manner. The crack can do this at various points and spread into the new grain. Such a

process gives rise to the formation of a number of steps that can group together, generating

a river marking. The convergence of tributaries is always in the direction of flow of the river

(i.e., "downstream"). This fact furnishes the possibility of determining the local direction of

propagation of crack in a micrograph. Under normal circumstances, face-centred cubic (FCC)

metals do not show cleavage. In these metals, a large amount of plastic deformation will

occur before the stress necessary for cleavage is reached. Cleavage is common in body-

centred cubic (BCC) and hexagonal close-packed (HCP) structures, particularly in iron and

low-carbon steels (BCC). Tungsten, molybdenum, and chromium (all BCC) and zinc

beryllium, and magnesium (all HCP) are other examples of metals that commonly show

cleavage.

Quasi cleavage is a type of fracture that is formed when cleavage occurs on a very fine scale

and on cleavage planes that are not very well defined. Fine grain sizes and higher

temperatures can lead to the occurrence of quasi-cleavage, which blends cleavage facets

with areas of dimple (MVC) rupture, such that the cleavage steps become tear ridges.

Typically, one sees this type of fracture in quenched and tempered steels. These steels

contain tempered martensitic and a network of carbide particles whose size and distribution

can lead to a poor definition of cleavage planes in the austenite grain. Thus, the real

cleavage planes are exchanged for small and ill-defined cleavage facets that initiate at the

carbide particles. Such small facets can give the appearance of a much more ductile fracture

than that of normal cleavage, and generally, river markings are not observed.




Intergranular Fracture

Intergranular fractures are typically the result of sustained load fractured, or a lack of

ductility in the material owing to segregation of embrittling elements and particles and

precipitates to the grain boundaries. For instance in temper embrittled steels and overaged

Al-Zn-Mg-Cu aluminium alloys. It is not possible to distinguish macroscopically between

intergranular fracture and brittle transgranular fracture: both appear faceted. However,

metallographic cross-sections through fracture surfaces and cracks will show whether the

fracture path is intergranular.

There are two main types of intergranular fracture appearance:

1) Grain boundary separation with microvoid coalescence. This type of intergranular

fracture occurs during overload failure of some steels and aluminium alloys, and also other

materials.

2) Grain boundary separation without microvoid coalescence. This type of intergranular

fracture occurs during overload failure of temper-embrittled steels and refractory metals

like tungsten, and also during sustained load fracture (creep, stress corrosion cracking,

embrittlement by hydrogen and liquid metals).

The dimples on the grain boundary facets are the main distinguishing feature of

intergranular fracture with microvoid coalescence. Intergranular fractures are not always

readily identifiable. Figure shows schematically an intergranular fracture along flat

elongated grains, which often occur in rolled sheet and plate materials as a consequence of

mechanical working. This type of intergranular fracture exhibits few grain boundary

junctions and is relatively featureless

Brittle intergranular fracture without MVC Brittle intergranular fracture with MVC




Ductile to Brittle transition

Under conditions of low temperatures, rapid loading and/or high constraint (e.g. when the

principal stresses are essentially equal), even ductile material may not exhibit any

deformation before failure. In ductile to brittle transition region fracture is controlled by the

competition between ductile tearing and cleavage fracture. The ductile crack growth occurs

by void growth and coalescence process which is driven by the increasing strain and

cleavage fracture occurs by a stress controlled process. Since a high crack tip can promote

cleavage fracture conditions and low constraint can promote ductile void growth

mechanism at temperatures in the mid ductile to brittle transition regime, either type of

fracture can occur depending on chosen geometry of size of the specimen.

At low temperatures, steel is brittle and fails by cleavage. At high temperatures, material is

ductile and fails by microvoid convalescence. Ductile fracture initiates at a particular

toughness value, the crack grows as load is increased and the specimen fails by plastic

collapse or tearing instability. In transition region between ductile and brittle behaviour,

both mechanisms of fracture can happen in the same specimen. In lower transition region

the fracture is by cleavage, but the toughness increases rapidly with temperature as

cleavage becomes more difficult.

Three basic factors contribute to a brittle-cleavage type of fracture. They are

1. a triaxial state of stress,

2. a low temperature, and

3. a high strain rate or rapid rate of loading.

All three of these factors do not have to be present at the same time to produce brittle

fracture. A triaxial state of stress, such as exists at a notch, and low temperature are

responsible for most service failures of the brittle type. However, since these effects are

accentuated at a high rate of loading, many types of impact tests have been used to

determine the susceptibility of materials to brittle behaviour. Steels which have identical

properties when tested in tension or torsion at slow strain rates can show pronounced

differences in their tendency for brittle fracture when tested in a notched-impact test.




Notched-bar Impact Tests

Various types of notched-bar impact tests are used to determine the tendency of a material

to behave in a brittle manner. This type of test will detect differences between materials

which are not observable in a tension test. The results obtained from notched-bar tests are

not readily expressed in terms of design requirements, since it is not possible to measure

the components of the triaxial stress condition at the notch. Furthermore, there is no

general agreement on the interpretation or significance of results obtained with this type of

test.

A large number of notched-bar test specimens of different design have been used by

investigators of the brittle fracture of metals. Two classes of specimens have been

standardized for notched-impact testing. Charpy bar specimens are used most commonly in

the United States, while the Izod specimen is favored in Great Britain.

The Charpy specimen has a square cross section (10x10 mm) and contains a 45° V notch, 2

mm deep with a 0.25 mm root radius. The specimen is supported as a beam in a horizontal

position and loaded behind the notch by the impact of a heavy swinging pendulum. The

specimen is forced to bend and fracture at a high strain rate of the order 103 s

-1. The Izod

specimen, which is used rarely today, has either a circular or square cross section and

contains a V notch near the clamped end.

The principal measurement from the impact test is the energy absorbed in fracturing the

specimen. After breaking the test bar, the pendulum rebounds to a height which decreases

as the energy absorbed in fracture increases. The energy absorbed in fracture, usually

expressed in joules, is rending directly from a calibrated dial on the impact tester.

The notched-bar impact test is most meaningful when conducted over a range of

temperatures so that the temperature at which the ductile-to-brittle transition takes place

can be determined.

The principal advantage of the Charpy V-notch impact test is that it is a relatively simple test

that utilizes a relatively cheap, small test specimen. Tests can readily be carried out over a

range of sub ambient temperatures. Moreover, the design of the test specimen is well

suited for measuring differences in notch toughness in low-strength materials such as

structural steels. The test is used for comparing the influence of alloy studies and heat

treatment on notch toughness. It frequently is used for quality control and material

acceptance purposes.







Ductile to Brittle Transition-Temperature Curve (DBTT)

The chief engineering use of the Charpy test is in selecting materials which are resistant to

brittle fracture by means of transition-temperature curves. The absorbed energy is plotted

against the testing temperature curve to give the Ductile to Brittle Transition-Temperature

Curve (DBTT). The curve represents change in behaviour from ductile at high temperature to

brittle at lower temperature. The design philosophy is to select a material which has

sufficient notch toughness when subjected to severe service conditions so that the load-

carrying ability of the structural member can be calculated by standard strength of materials

methods without considering the fracture properties of the material or stress concentration

effects of cracks or flaws.

Lower shelf transition mixed mode




Upper shelf

The design philosophy using transition temperature curve centres about the determination

if temperature above which brittle fracture will not occur at elastic stress levels. Obviously,

lower this transition temperature, the greater the fracture toughness of the material




Criterion for Transition Temperature

The shape of a typical energy absorbed to temperature curve shows that there is no single

criterion that defines the transition temperature. The various definitions of transition

temperature obtained from energy vs. temperature curve or fracture appearance vs.

temperature curve are illustrated below.

The most comprehensive criterion for transition temperature is to select T1, corresponding

to the upper shelf in the fracture energy and the temperature above which the fracture is

100% fibrous. This transition temperature criterion is called fracture transition plastic (FTP).

The FTP is the temperature at which the fracture changes from totally ductile to

substantially brittle. The probability of brittle fracture is negligible above the FTP. The use of

FTP is conservative and in many applications impractical.

An arbitrary but less conservative criterion is to base the transition temperature on 50%

cleavage -50% shear, T2. This is called fracture appearance transition temperature (FATT).

Correlations between Charpy impact test and service failure indicate that less than 70%

cleavage fracture in Charpy bar indicates a high probability that the failure will not occur at

or above the temperature is the stress does not exceed about one half of the yield stress.

Roughly similar results are obtained by defining the transition temperature as the average if

the upper and lower shelf values, T3.

A common criterion is to define the transition temperature, T4 on the basis of an arbitrary

low value of energy absorbed. This is often called ductility transition temperature. For low

strength ships this value is taken 20J. But for other material this value is not known.

A well defined criterion is to base the transition temperature on the temperature at which

the fracture becomes 100% cleavage, T5. This point is known as nil ductility temperature

(NDT). The NDT is the temperature at which fracture initiates with essentially no prior

plastic deformation. Below the NDT probability of ductile fracture is negligible.

Various criteria of transition temperature




Metallurgical Factors affecting Transition Temperature.

The shape and position of DBTT curve is important as it determines the transition

temperature, which indicates where it is safe to use for the given application. There are

several factors affecting DBTT curve.

• Crystal structure

• Interstitial atom

• Grain size

• Heat treatment

• Specimen orientation

• Specimen thickness

Effect of Crystalline Structure

The transition-temperature behaviour of a wide spectrum of materials falls into the three

categories. Medium- and low-strength fcc metals and most hcp metals have such high notch

toughness that brittle fracture is not a problem unless there is some special reactive

chemical environment. High-strength materials (s0 > E/150) have such low notch toughness

that brittle fracture can occur at nominal stresses in the elastic range at all temperatures

and strain rates when flaws are present.

High-strength steel, aluminum and titanium alloys fall into this category. At low temperature

fracture occurs by brittle cleavage, while at higher temperatures fracture occurs by low-

energy rupture. It is under these conditions that fracture mechanics analysis is useful and

appropriate.

The notch toughness of low- and medium-strength bcc metals, as well as Be, Zn, and

ceramic materials is strongly dependent on temperature. At low temperature the fracture

occurs by cleavage while at high temperature the fracture occurs by ductile rupture. Thus,

there is a transition from notch brittle to notch tough behaviour with increasing

temperature. In metals this transition occurs at 0.1 to 0.2 of the absolute melting

temperature Tm, while in ceramics the transition occurs at about 0.5 to 0.7 Tm.

Effect of Crystalline Structure on Transition temperature




Effect of Interstitial atom

Changes in transition temperature of over 55°C (100°F) can be produced by changes in the

chemical composition or microstructure of mild steel. The largest changes in transition

temperature result from changes in the amount of carbon and manganese. This transition

temperature is lowered about 5.5°C (10°F) for each increase of 0.1 percent manganese and

raised by about 14oC for each increase of 0.1% carbon. Increasing the carbon content also

has a pronounced effect on the maximum energy and the shape of the energy transition-

temperature curves.

The Mn/C ratio should be at least 3/1 for satisfactory notch toughness. A maximum

decrease of about 55°C (100°F) in transition temperature appears possible by going to

higher Mn/C ratios.

Phosphorus also has a strong effect in raising the transition temperature. The temperature

is increased by 7oC for every .01% phosphorous. The role of nitrogen is difficult to assess

because of its interaction with other elements. It is, however, generally considered to be

detrimental to notch toughness.

Nickel is generally accepted to be beneficial to notch toughness in amounts up to 2 percent

and seems to be particularly effective in lowering the ductility transition temperature.

Silicon, in amounts over 0.25 percent, appears to raise the transition temperature.

Molybdenum raises the transition almost as rapidly as carbon, while chromium has little

effect.

Notch toughness is particularly influenced by oxygen. When oxygen content was raised from

.001% to .053% the transition temperature was raised from -15oC to 340

oC

Effect of carbon content in the energy-transition-temperature curves for steel




Effect of Grain Size

Grain size has strong effect on transition temperature. An increase of one ASTM number in

the ferrite grain size (actually a decrease in grain diameter) can result in decrease in

transition temperature of 16oC for mild steel. Decreasing grain diameter from ASTM grain

size 5 to ASTM size 10 can change the transition temperature from about 20oC to -50

oC. A

similar effect of decreasing with transition temperature with decreasing austenitic grain size

is observed with higher alloyed heat treated steels. Many of the variables concerned with

processing mild steel affect the ferrite grain size and therefore affect the transition

temperature. Air cooling and aluminium oxidisation results in lower transition temperature.

Using lowest possible finishing temperature for hot rolling is beneficial. Spray cooling from

rolling temperature before coiling can lower the transition temperature by 500C.

Effect of grain size on DBTT




Effect of Heat-Treatment

Low carbon steel can exhibit two types of aging phenomenon which produce an increase in

transition temperature. Quench aging is caused by carbide precipitation in a low carbon

steel which has been quenched from around 700oC. Strain aging occurs in low carbon steel

which has been cold worked. Cold working by itself will increase the transition temperature,

but strain aging results in greater increase, usually around 25oC to 30

oC.

Tempered martensitic structure produces the best combination of strength and impact

resistance of any microstructure that can be produced in steel. The tensile properties of

tempered martensitic of the same hardness and carbon content are alike, irrespective of the

amount of other alloy additions. This generalization holds approximately for room

temperature. Impact resistance of heat treated steels, but it is not valid for the variation of

impact resistance with temperature. Figure shows the temperature dependence for impact

resistance for a number of different alloy steels, all having about .4% Carbon and all

tempered martensitic structure produced by quenching and tempering to a hardness of

Rc35. Note that the maximum variation of about 100oC in the transition temperature at 30 J

level is possible. Every greater spread in transition temperature would be obtained if the

tempering temperature were adjusted to give a higher hardness.

Temperature dependence of impact resistance for different alloy steel of same

carbon content quenched and tempered to Rc35




Effect of Specimen Orientation

The notched impact properties of rolled or forged products vary with the orientation in the

plate or bar. The figure shows typical form of energy-temperature curves for specimen cit in

longitudinal and transverse directions of the rolled plate. Specimen A and B are oriented in

longitudinal directions. The graphs shows that considerably large differences are expected

for different specimen orientations at high energy levels, but difference becomes much less

at energy levels below 30J. Since the ductility transition temperatures are evaluated in this

region of energy, it seems that specimen and notch orientation are not very important. If,

however materials are compared on the basis of room temperature impact properties,

orientation can greatly affect the results.

Effect of orientation of specimen on transition temperature




Effect of specimen thickness

Probably the chief deficiency of Charpy impact test is that the small specimen is not always

a realistic model of the actual situation. Not only does the small specimen lead to

considerable scatter, but a specimen with a thickness of 10mm cannot provide the same

constraint as would be found in a structure with a much greater thickness, at a particular

service temperature the standard Charpy specimen shows a high shelf energy, while actually

the same material in a thick section structure has low toughness at the same temperature.

Effect of section thickness on transition-temperature curves




Conclusion

In most design situations a material that demonstrates ductile fracture is usually preferred

for several reasons.

First and foremost, brittle fracture occurs very rapidly and catastrophically without any

warning. Ductile materials plastically deform, thereby slowing the process of fracture and

giving ample time for the problem to be corrected.

Second, because of the plastic deformation, more strain energy is needed to cause ductile

fracture.

Next, ductile materials are considered to be "forgiving" materials, because of their

toughness you can make a mistake in the use, design of a ductile material and still the

material will probably not fail. Also, the properties of a ductile material can be enhanced

through the use of one of the strengthening mechanisms. Strain hardening is a perfect

example, as the ductile material is deformed more and more its strength and hardness

increase because of the generation of more and more dislocations.

Therefore, in engineering applications, especially those that have safety concerns involved,

ductile materials are the obvious choice. Safety and dependability are the main concerns in

material design, but in order to attain these goals there has to be a thorough understanding

of fracture, both brittle and ductile. Understanding fracture and failure of materials will lead

the materials engineer to develop safer and more dependable materials and products.




References

1. Fracture mechanics by Michael Janssen, Jan Suidema, Russel Wanhill.

2. Mechanical Behaviour of Material by Marc Andre Meyers, Krishnan Kumar Chawla.

3. Fracture Mechanic – Fundamental & Application by T. L. Anderson

4. Non Linear Fracture Mechanics for Engineers by Ashok Saxena

5. Mechanical Metallurgy by George.E.Dieter

Documents

Seminar on Brittle and Ductile Fracture