Seminar on Brittle and Ductile Fracture

2008

Seminar on Brittle and Ductile Fracture

Department of Mechanical Engineering

Manipal Institute of Technology

Chetan Purushottam Bhat Mtech (CAMDA)


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

Manipal Institute of Technology Department of Mechanical Engineering


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: 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 brittle and ductile material

pure ductile

moderately ductile

fracture


fracture

brittle 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 pu rity 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 ide al 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 2 0 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 AI 2O3, in aluminium and ThO 2, in nickel. Precipitate particles obtained by appropriate heat treatment also form part -of this class (eg., an Al-cu-Mg Manipal Institute of Technology Department of Mechanical Engineering


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 ve ry 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

,

and and

are are the the prin princi cipa pall norm normal al stres stresse ses. s. Acco Accord rdin ing g to the the Argo Argon, n, et al model model,, the 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 o penny-shaped flaw produces deformation bands at 45 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 o the latter surface is oriented 45 from the tensile axis and there is little evidence (at low magnifications) of microvoid coalescence, many refer to this type of surface as shear o fracture. The 45 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 e ach 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 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 n umber 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 bodycentred 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 b rittle 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 b rittle type. However, since these effects are accentuated at a high rate of loading, many types o f 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 d esign 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 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 3 -1 specimen is forced to bend and fracture at a high strain rate of the order 10 s . The Izod specimen, which is used rarely today, has either e ither a circular or square cross section and contains a V notch near the clamped end. The principal measurement from the impact test is the ene rgy 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 ene rgy 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 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 loadcarrying 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 T 1, 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, T 4 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, T 5. 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 (s 0 > 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 lowenergy 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 T m.

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 o raised by about 14 C 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 transitiontemperature 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 o is increased by 7 C 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. Whe n oxygen content was raised from o o .001% to .053% the transition temperature was raised from -15 C to 340 C

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 o transition temperature of 16 C for mild steel. Decreasing grain diameter from ASTM grain o o size 5 to ASTM size 10 can change the transition temperature from about 20 C to -50 C. 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 0 rolling temperature before coiling can lower the transition temperature by 50 C.

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 o

steel which has been quenched from around 700 C. Strain aging occurs in low carbon steel which has been cold worked. Cold working by itself will increase the transition temperature, o

o

but strain aging results in greater increase, usually around 25 C to 30 C. 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 o

Rc35. Note that the maximum variation of about 100 C 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 R c35



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 see ms that specimen and notch orientation are not very important. If, however materials are compared on the basis b asis 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


Seminar on Brittle and Ductile Fracture

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